TIMIS2012 141B t Annual Meeting & Exhibition

S u p p l e m e n t a l P roceed ings Volume 1:

Materials Processing and Interfaces

TIMIS2012 141 s t Annual Meeting & Exhibition

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Energy Technology 2012: C02 Management and Other Technologies

EPD Congress 2012

International Smelting Technology Symposium (Incorporating the 6th Advances in Sulfide Smelting Symposium)

Light Metals 2012

Magnesium Technology 2012

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces

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TIMS2012 141 s t Annual Meeting & Exhibition

Supplemental Proceedings Volume 1:

Materials Processing and Interfaces

About this volume

The TMS 2012 Annual Meeting Supplemental Pmceeclings, Volume I: Materials Pir>cessinx and Interfaces, is a collection of papers from the 2012 TMS Annual Meeting and Exhibition, held Marchl l-March 15, in Orlando. California, U.S.A.

The papers in this volume were selected based on technical topic compatibility and represent ten symposia from the meeting. This volume, along with the other proceedings volumes published for the mecting,.and archival journals, such as Metallurgical anil Materials Transactions and lite Journal of Electronic Materials, represents the available written record of the 65 symposia held at the 2012 TMS Annual Meeting.

The individual papers presented within this proceedings volume have nut necessarily been edited or reviewed by the conference program organizers and are presented ''as is." The opinions and statements expressed within the papers are those of the individual authors only and are not necessarily those of anyone else associated with the proceedings volume, the source conference, or TMS. No confirmations or endorsements are intended or implied.

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Copyright © 2012 by The Minerals, Metals, & Materials Society. All rights reserved.

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TABLE OF CONTENTS Supplemental Proceedings: Volume I: Materials Processing and Interfaces

Advances in Surface Engineering: Alloyed and Composite Coatings

Session II

Structural Coatings in Aluminum Alloy Microtruss Materials 3 B. Yu, andG. Hibbard

Laser Cladding of High-Performance CPM Tool Steels on Hardened HI 3 Hot-Work Tool Steel for Automotive Tooling Applications 11

J. Chen, andL. Xue

Electron Beam Deposited Multilayer Optical Interference Coatings Using Oxide Composites 19

A. Nayak, N. Sahoo, R. Tokas, A. Biswas, andN. Kamble

Session III

Microstructure and Wear Properties of Laser In-situ Formation of TiBx and TiC Titanium Composite Coatings 27

J. Liang, C. Liu, S. Chen, and C. Ren

Creep Behavior of Plasma Sprayed Y-PSZ Coated 6063-T6 Aluminum Alloy 35

E. Erzi, C. Kahruman, andS. Yilmaz

Contribution of Ti Addition to the Electronic Structure and Adhesion at the Fe2Al5/Fe Interface in 55%A1-Zn Coating 41

G. Wu, Y. Ren, J. Zhang, andK. Chou

v

Session IV

The Roles of Diffusion Factors in Electrochemical Corrosion of TiN and CrN (CrSiCN) Coated Mild Steel and Stainless Steel 49

F. Cai, Q. Yang, andX. Huang

Effect of Electroplating Parameters on "HER" Current Density in Ni-MoS2 Composite Plating 57

E. Saraloglu Guler, I. Karakaya, andE. Konca

Production of Ceramic Layers on Aluminum Alloys by Plasma Electrolytic Oxidation in Alkaline Silicate Electrolytes 65

A. Lugovskoy, A. Kossenko, B. Kazanski, andM. Zinigrad

Abrasive Wear Properties of Plasma Sprayed Y-PSZ Coated 6063-T6 Aluminum Alloy 73

E. Erzi, S. Yildirim, andS. Yilmaz

The Electrochemical Behavior of Surgical Grade 316L S S with and without HA Coatings in Simulated Body Fluid 79

T. Singh, H. Singh, H. Singh, andH. Saheet

Modification Research on the Influence on Corrosion Film Properties of Pb-Ca-Sn Alloys of with Addition of Bi, Ag and Zn 87

L. Xu, L. Liu, and P. Zhu

Session V

Evaluation of Residual Stress in Fe2B Coating on Ductile Cast Iron 95 M. Donu Ruiz, N. Lopez Perrusquia, V. Cortez Suarez, and D. Sanchez Huitron

General Poster Session

Session I

Influence of Heat Treatment on the Corrosion of Steels in CCS Environment 103

A. Pfennig, S. Schulz, A. Kranzmann, T. Werlitz, S. Wetzlich, E. Billow, J. Tietböhl, and C. Frieslich

vi

Microstructure and Property Modifications in Mould Steels Treated by Pulsed Electron Beam I l l

K. Zhang

Potential Fiberboard Material from Cow Manure and Disposable Water Bottle 119

B. Ng, M. Murray, C. Bradfield, andR Pritish

Influence of Process and Thermo-physical Parameters on. the Heat Transfer at Electron Beam Melting of Cu and Ta 125

K. Vutova, V. Donchev, V. Vassileva, and G. Mladenov

Industrial Use of a New Ultrasound Spray for Cooling and Wet Gas Treatment in the Pyrometallurgical Processes 133

M. Cirkovic, V. Trujic, and Z. Kamberovic

Development of 3D Porous Nickel Electrodes for Hydrogen Production 141 V. Pérez-Herranz, I. Herrâiz-Cardona, E. Ortega, andJ. Garcia-Anton

Electrochemical Recovery of Zinc Present in the Spent Pickling Baths Coming from Hot Dip Galvanizing Processes 149

V. Pérez-Herranz, J. Carrillo-Abad, M. Garcia-Gabaldôn, andE. Ortega

Laboratory Testing Results of Kinetics and Processing Technology of the Polymetallic Sulphide Concentrate Blagojev Kamen - Serbia 157

M. Cirkovic, Z. Kamberovic, and V. Trujic

Hidrotalcite with Gentamicine, of the Type Mg0.68Alo.32(OH)2

(N03)o.32*0.1H20, Formed by Chemical Coprecipitation in Controlled Atmosphere 165

H. Rodriguez-Santoyo, and O. Martinez-Alvarez

Effect of Thiodiglycolamide Addition to Di-n-hexyl Sulfide on the Pd(II) Extraction Rate 173

H. Narita, M. Tanaka, andS. Ueno

Synthesis and Characterization of Metallic Oxides 179 E. Brocchi, R. Souza, M. Doneda, J. Campos, A. Wimmer, andR. Navarro

Fabrication of Lotus-Type Porous Copper by Centrifugal Casting Technique 187

Y. Lee, H. Kim, M. Kim, andS. Hyun

vii

Effects of Pulsed Magnetic Annealing on the Grain Boundary of Primary Recrystallized Microstructure in the Grain-Oriented Silicon Steel 191

J. Huang, L. Liu, X. Xia, X. Jiang, L. Li, and Q. Zhai

Relationship between Heat Input and Microstructure and Mechanical Properties of Laser Beam Welded Superalloy Inconel 718 199

A. Odabasi, N. Ünlti, G. Göller, M. Eruslu, andE. Kayali

Experimental Study on the Behavior of Slag Entrapment and Inclusion Removal in 441 Ladle with Argon Blowing 207

S. Zheng, M. Zhu, andZ. Cheng

The Effect of High Superheat on the Solidification Structure and Carbon Segregation of Ferrite-Based Alloy 215

H. Zhong, Y. Tan, H. Li, X. Mao, and Q. Zhai

Refinement of Ligaments of Nanoporous Ag Ribbons by Controlling the Surface Diffusion of Ag 223

T. Song, Y. Gao, Z. Zhang, and Q. Zhai

Thermal Analysis of the Composition of Poly(Acrylic AcidyCarboxymethylstarch Used as a Polymeric Binder 231

B. Grabowska, M. Holtzer, S. Eichholz, K. Hodor, andE. Olejnik

Mechanical Behavior Related to Interface Physics

Grain Boundaries; Experiment and Modeling

Interfacial Strength of Al/Zr/DU-10%wtMo Subject to Different Loading Modes 241

M Lovato, C. Liu, and W. Blumenthal

Interface Evolution under Mechanical Loading: Experiment, Characterization, and Theoretical

Modeling

Uniaxial Tension of Friction-Welded 304-Stainless Steel and 6061 Aluminum 249

C. Liu, M. Lovato, and W. Blumenthal

via

Interface Structures; Characterization, Theory, and Modeling

Ultra Fast Grain Boundary Segregation In Hot Deformed Nickel 257 M Allart, F. Christien, and R. Le Gall

Quantitative NanoSIMS Analysis of Grain Boundary Segregation in Bulk Samples 265

F. Christien, K. Moore, C. Downing, andC. Grovenor

The Periodic Unit of Doubly-diffracted Reflections from Periodic Grain Boundaries in Cubic Crystals and Its Relationship with Coincident Site Lattice 273

M. Shamsuzzoha

Poster Session

The In-situ Intrinsic Stress Measurements of Cu and Al Thin Films 281 J. Yu, and Y. Kim

Delamination Characterization of Bonded Interface Using Surface Based Cohesive Model 289

M. Ramamurthi, and Y. Kim

Nanocomposites

Mechanical Behavior and Modelling of Nanocomposites

Compressive Strength of Epoxy- Graphite Nanoplatelets Composites 299 H. Colorado, A. Wong, andJ. Yang

Micromechanical Analysis of Influences of Agglomerated Nanotube Interphase on Effective Material Properties of a Three Phase Piezoelectric Nanocomposite 307

T. Tang, and P. Wang

Effect of Nano-Paper Coating on Flexural Properties of a Fire-Treated Glass Fiber-Reinforced Polyester Composite 313

J. Skovron, J. Zhuge, A. Gordon, J. Kapat, andJ. Gou

IX

Finite Element Modeling of the Nanoscratching of Polymer Surfaces 321 W. Chirdon, andJ. Rozas

Processing of Nanocomposites I

Manufacturing and Characterization of an Auxetic Composite 329 F. Chiang

Microtruss Cellular Nanocomposites 337 K. Abu Samk, G. Huang, M. Skocic, H. Zur ob, D. Embury, O. Bouaziz, and G. Hibbard

Nanocomposites for Magnetic and Dielectric Applications

Synthesis of Tailored Core-Shell Magnetic Microparticles for Intravascular Embolization 345

G. Ferreira, A. Umpierre, andF. Machado

Dramatic Expansion of Luminescence Region in GaP/Polymer Nanocomposites 353

S. Pyshkin, andJ. Ballato

Nanocomposite Interfaces and Characterization

Positron Lifetime Analysis of Polyurea-Nanoclay Composites 361 N. Seetala, D. Hubbard, G. Burks, A. Trochez, and V. Khabashesku

Processing of Nanocomposites II

Rheological Properties of Suspensions of Nanopowders Yttrium Oxide (Y203) and Magnesium-Aluminum Spinel (MgAl204) 367

G. Zyla, M. Cholewa, A. Witek, J. Plog, V. Lehmann, T. Oerther, andD. Gross

Thermal Properties of Hemp-High Density Polyethylene Composites: Effect of Two Different Chemical Treatments 375

N. Lu, and S. Oza

x

Discarded Ultrafine Particulate Carbonaceous Materials Used as Reinforcers of Rubber Vulcanized Products 383

G. Martin-Cortés, F. Esper, L. Galvào Dantas, W. Hennies, andF. Valenzuela-Diaz

Properties of Additional Reinforcers Materials Used to Complement NAOB -A Rubber / Organoclay Nanocomposite Material 389

F. Esper, G. Martin-Cortés, L. Dantas, A. Cutrim, W. Hennies, andF. Valenzuela-Diaz

Poster Session

Thermal Properties of Carbon Nano Tubes Reinforced Mg-Matrix Nanocomposites 395

S. Iqbal, A. Mustafa, S. Talapatra, and P. Filip

New Advances in Synthesis, Characterization, and Application of Layered Double Hydroxides

Session I

Designing Layered Double Hydroxides for Targeted Applications 405 J. Hossenlopp, S. Majoni, and C. Machingauta

Electrochemical Synthesis of Layer Double Hydroxides, Its Characterization, and Performance Study for Removal of Nitrate and Arsenic 413

M. Haider, J. Gomes, K. Urbanczyk, D. Cocke, H. McWhinney, G. Irwin, and P. Bernazzani

Removal of Direct Red and Orange II Azo Dye from Synthetic Textile Water Using Electrochemically Produced Fe-LDH 421

S. Jame, J. Gomes, and D. Cocke

Removal of Arsenic Using Green Rust and Other Electrochemically Generated Floe 429

M. Rahman, J. Gomes, K. Urbanczyk, andD. Cocke

Formation of Layered Double Hydroxides in Self-Purification of Polynary Metal Electroplating Wastewaters for Effective Removal of Anionic Dye 437

J. Zhou, G. Qian, C. Liu, Y. Wu, X. Ruan, Y. Xu, andJ. Liu

XI

Characterization and Chemical Modification of Electrochemically Produced Layered Double Hydroxides as Nanomaterials 445

M. Islam, J. Gomes, and P. Bernazzani

Randall M. German Honorary Symposium on Sintering and Powder-Based Materials

Sintering Theory and Practice

A Review on Alloying in Tungsten Heavy Alloys 455 A. Bose, R. Sadangi, and R. German

Current Activated and Conventional Sintering

Low-Thermal Load Consolidation of Sm-Fe-N Flake Powder by Combination of Cyclic Compression and Current Sintering 467

K. Takagi, H. Nakayama, andK. Ozaki

Fabrication of TiN / Fe-Al Cermet from Mixture of TiN, Fe and Al Powders 475

H. Nakayama, K. Ozaki, andK. Kobayashi

Transparent Polycrystalline Alumina Obtained by SPS: Single and Double Doping Effect 481

B. Apak, H. Kanbur, E. Ozkan Zayim, G. Goller, O. Yucel, andF. Cinar Sahin

Sintering of Nanocrystalline Tungsten Powder 489 W. de Rosset

Powder Technology

Effect of Powder Synthesis and Processing on Luminescence Properties 497 J. McKittrick, J. Han, J. Choi, andJ. Talbot

Effect of Rapid Solidification and Heat Treatment on D2 Tool Steel 505 P. Delshad Khatibi, H. Henein, andD. Ivey

XI1

Powder Processing and Consolidation I

Development of Solid Freeform Fabrication for Metallic Parts Using Selective Inhibition of Sintering 513

M. Yoozbashizadeh, andB. Khoshnevis

Numerical Simulation of Cold Pressing of Armstrong CP-Ti Powders 521 A. Sabau, S. Gorti, W. Peter, W. Chen, and Y. Yamamoto

The Effect of Coke Particle Size on the Thermal Profile of the Sintering Process Product 529

N. Tahanpesarandezfuly, and A. Heidary Moghadam

Powder Processing and Consolidation II

Powder Material Principles Applied to Additive Manufacturing 537 D. Bourell, andJ. Beaman

Processing Challenges of Dual-Matrix Carbon Nanotube Aluminum Composites 545

A. Esawi, K. Morsi, I. Salama, and H. Saleeb

Influence of High Pressure Torsion on the Consolidation Behavior and Mechanical Properties of AA6061-SiCp Composite Powders 553

H. Salem, W. El-Garaihy, andE. Rassoul

Powder Processing and Consolidation III

LASER Powder Deposition of AlMgBi4-TiB2 Ultra-Hard Coatings on Titanium and Steel Substrates 561

J. Fuerst, M. Carter, andJ. Sears

Mechanical Properties of Spark Plasma Sintered ZrC-SiC Composites 569 S. Sagdic, I. Akin, F. Sahin, O. Yucel, and G. Goller

Intense Pulsed Light Sintering Technique for Nanomaterials 577 H. Colorado, S. Dhage, J. Yang, andH. Hahn

xni

Recent Developments in Biological, Electronic, Functional and Structural Thin Films and Coatings

Process-Properties-Performance Correlations I

Dependence of Tribology of Carbide Derived Carbon Films on Humidity 587 M. Tlustochowicz

Structural and Optical Properties of Silicon Carbonitride Thin Films Deposited by Reactive DC Magnetron Sputtering 595

O. Agirseven, T. Tavsanoglu, E. Ozkan Zayim, and O. Yucel

Influence of TIG Re-Melting and RE (La203) Addition on Microstructure, Hardness and Wear of Ni-WC Composite Coating 603

B. Dhakar, D. Dwivedi, andS. Sharma

Evaluation of Mechanical Properties of Ni-Ti Bi-Layer Thin Film 609 M. Mohri, and M. Nili-Ahmadabadi

Anodic Ti02 Nanotubular Arrays with Pre-Synthesized Hydroxyapatite - A Promising Approach to Enhance the Biocompatibility of Titanium 617

L. Wang

Preparation and Properties of Cu2ZnSnS4 Thin Films by Electrodeposition and Sulfurization 625

C. An, H. Lu, andX. Chen

Process-Properties-Performance Correlations II

Formation of Crystalline and Amorphous Phases During Deposition of NixTi].x

Thin Film on Si Substrate - Interpretation of Experimental Results Using Molecular Dynamics Simulations 633

S. A ich, G. Priyadarshini, M. Gupta, S. Ghosh, and M. Chakraborty

Applications to Bio, Energy and Electronic Systems

Doping and Co-Doping of Bandgap-Engineered ZnO Films for Solar Driven Hydrogen Production 641

S. Shet, N. Ravindra, Y. Y an, and M. Al-Jassim

Magnetic Field Assisted Heterogeneous Device Assembly 651 V. Kasisomayajula, M. Booty, A. Fiory, andN. Ravindra

xiv

Organic Thermal Mode Photoresists for Applications in Nano-Lithography ...663 H. Wu, M. Li, C. Yang, C Cheng, S. Chen, andD. Huang

Process-Properties-Performance Correlations III

Nitrogen Doped ZnO (ZnO:N) Thin Films Deposited by Reactive RF Magnetron Sputtering for PEC Application 669

S. Shet, K. Ahn, N Ravindra, Y. Y an, and M. Al-Jassim

Spin-Coated Erbium-Doped Silica Sol-Gel Films on Silicon 677 S. Abedrabbo, B. Lahlouh, S. Shet, A. Fiory, andN. Ravindra

Influence of Annealing on the Martensitic Transformation and Magnetocaloric Effect in Ni49Mn39Sni2 Ribbons 685

D. Wu, S. Xue, H. Zheng, andQ. Zhai

Metal Diaphragm Based Magnetic Field Sensor 693 A. Banobre, I. Padron, A. Fiory, andN. Ravindra

Optical and Electronic Properties of AlN, GaN and InN: An Analysis 701 C. Lamsal, D. Chen, andN. Ravindra

Science and Engineering of Light Metal Matrix Nanocomposites and Composites

Metal Matrix Nanocomposites

Interfacial Analysis of CNT Reinforced AZ61 Mg Alloy Composites 717 K. Kondoh, H. Fukuda, J. Umeda, and B. Fugetsu

Biodegradability and Mechanical Performance of Hydroxyapatite Reinforced Magnesium Matrix Nanocomposite 725

C. Ma, L. Chen, J. Xu, A. Fehrenbacher, Y. Li, F. Pfefferkorn, N. Duffie, J. Zheng, andX. Li

Mechanical Properties of A356-CNTCast Nano Composite Produced by a Special Compocasting Route 733

B. Abbasipour, B. Niroumand, andS. Monirvaghefi

Production of Cast AZ91-CNT Nano-Composite by Addition of Ni-P-CNT Coated Magnesium Powder to the Melt 741

M. Firoozbakht, B. Niroumand, andS. Monirvaghefi

xv

Wear Behavior of Magnesium Matrix Nanocomposites at Room and Elevated Temperature 749

W. Li, andS. Liu

Nanocomposites and Composites

Uniform Dispersion of Nanoparticles in Metal Matrix Nanocomposites 757 L. Chen, H. Choi, A. Fehrenbacher, J. Xu, C. Ma, andX. Li

Effect of Core-shelled Nanoparticles of Carbon-Coated Nickel on Magnesium 765

Y. Sun, H. Choi, H. Konishi, V. Pikhovich, R. Hathaway, andX. Li

In Situ Composite of (Mg2Si)/Al Fabricated by Squeeze Casting 775 H. Lus, G. Ozer, andK. Guler

Optimization of Tensile Strength of Friction Stir Welded Al-(10 tol4 wt.%) TiB2 Metal Matrix Composites 783

S. Joseph Vijay, N. Murugan, andS. Parameswaran

Metal Matrix Composites

Slow-Shot High Pressure Die-Casting (SS-HPDC) Process for AE44 Magnesium Single-Cylinder Engine Block with Short-Fiber Reinforcement in the Bore 791

B. Hu, P. Wang, B. Powell, andX. Zeng

Compressive Properties of A1-B4C Composites over the Temperature Range of 25-500°C 799

S. Gangolu, A. Rao, N. Prabhu, V. Deshmukh, andB. Kashyap

Aluminum Metal Matrix Composite via Direct Metal Laser Deposition: Processing and Characterization 807

B. Wälder a, and S. Kalita

Damage Evolution Model for Hybrid Metal Matrix Composites 815 J. Dibelka, andS. Case

Numerical Simulation of Pressure Infiltration Process for Making Metal Matrix Composites: Effect of Process Parameters 823

B. Wang, andK. Pillai

xvi

Titanium: Advances in Processing, Characterization and Properties

Processing and Process Modeling I

Microstructural Evolution and Mechanical Properties of ß-Titanium Ti-lOV-2Fe-3Al during Incremental Forming 833

S. Winter, S. Fhtsch, andM. Wagner

Study on Hot Deformation Behavior of TC4 Titanium Alloy 841 Y. Lu, S. Jiao, X. Zhou, and A. Dong

Evolution of Microstructures and Properties of Ti-44Al-6V-3Nb-0.3Y Alloy after Forging and Rolling 849

Y. Chen, H. Niu, S. Xiao, P. Sun, and C. Zhang

Effect of Forging on Microstrutural Characteristic and Tensile Properties of In-Situ (TiB+TiC)/Ti Composite 857

Y. Chen, C. Zhang, S. Xiao, D. Wu, and H. Niu

Processing and Process Modeling II

Microstructure and Mechanical Properties of Ti-6A1-4V Fabricated by Selective Laser Melting 863

M. Simonelli, Y. Tse, and C. Tuck

Computational Modeling of the Dissolution of Alloying Elements 871 J. Ou, A. Chatterjee, D. Maijer, S. Cockcroft, and C. Reilly

Cost Effective and Eco-Friendly Process for Preparation of Wrought Pure Ti Material via Direct Consolidation of TiH2 Powders 879

T. Mimoto, N. Nakanishi, T. Threrujirapapong, J. Umeda, andK. Kondoh

Effect of Dual-Laser Beam Welding on Microstructure Properties of Thin-Walled y-TiAl Based Alloy Ti-45Al-5Nb-0.2C-0.2B (TNB) 887

J. Liu, V. Ventzke, P. Staron, H. Brokmeier, M. Oehring, N. Kashaev, andN. Huber

xvii

Microstructure Evolution and Characterization I

Recrystallization Behavior in Ti-13Cr-lFe-3Al Alloy after Severe Plastic Deformation 895

M. Ueda, H. Matsuhira, Y. Takasaki, M. Ikeda, and Y. Todaka

Mechanical Properties

Crack Initiation and Microstructurally Short Crack Growth of Ti-6A1-4V 903 H. Christ, H. Knobbe, P. Koester, C. Fritzen, and M. Riedler

Three-dimensional Investigation of the Microtexture near Tensile Crack Tip in Ti-6A1-4V 911

X. Xu, Y. Tse, G. West, and A. Huang

Machinability of ß-Titanium Alloy Ti-10V-2Fe-3Al with Different Microstructures 919

H. Abrahams, C. Mâchai, and D. Biermann

Residual Stress Relaxation Effects on the Cracking and Wear Processes of Shot Peened Ti-6A1-4V Titanium Alloy under Fretting-Fatigue Loading 927

R. Ferre, S. Fouvry, B. Berthel, R. Amargier, and A. Ferre

General Abstracts

Efficient Oxidation Protection of Ti- and TiAl-alloys by F-treatments 935 A. Donchev, M. Schütze, R. Yankov, and A. Kolitsch

Characteristics and Wear Performance of Nitrided Ti6A17Nb 941 F. Siyahjani, M. ïpekci, and H. Cimenoglu

Composition Analysis of Diffusion Bonded y-TiAl Intermetallic: TiAlV Alloy Interface by Using STEM 947

P. Sivagnanapalani, Gouthama, andM. Sujata

Effect of Erbium Addition on Microstructre of As-Cast Ti-22Al-25Nb Alloy 955

J. Dai, H. Lu, andZ. Cai

Fracture Behaviors of TiN and TiN/Ti Multilayer Coatings on Ti Substrate during Nanoindentation 963

Y. Sun, C. Lu, A. Tieu, Y. Zhao, H. Zhu, K. Cheng, and C. Kong

xvm

Deformation Mechanism in Nanoindentation of Ti63.375Fe34.i25Sn2 5 Alloy 971 K. Cheng, C. Lu, K. Tieu, L. Zhang, and Y. Sun

Wettability and Interfacial Phenomena between Metals and Ceramic/Refractory Materials

Session I

Chemical Wear of Basic Brick Linings in the Non-Ferrous Industry 981 D. Gregurek, A. Spanring, M Kirschen, andC. Majcenovic

Study on Wettability of Cu and 85Cu-15Ni Alloy on 18NiO-NiFe204

Composite Ceramics 989 J. Du, Y. Liu, G. Yao, Z. Zhang, and G. Zu

Interfacial Reactions in the Liquid/Solid and Liquid/Vapor Interfaces of Al-Si-Mg Alloys and B l 2 (Bc2) Substrates 997

O. Herr era-Romero, M. Pech-Canul, Z. Chaudhury, and G. Newaz

TMS2011 General Abstracts: Structural Materials Division

Microstructure

Microstructure of a' Martensites in Ti-V-Al Alloys Studied by High-Resolution Transmission Electron Microscopy 1007

K. Sato, H. Matsumoto, A. Chiba, and T. Konno

Author Index 1013

Subject Index 1019

xix

TIMIS2012 141 s t Annual Meeting & Exhibition

Advances in Surface Engineering: Alloyed and

Composite Coatings

The proceedings contained in this section have not been edited or reviewed by the conference program organizers. The opinions and statements expressed in the proceedings are those of the authors only and are not necessarily those of the editors or TMS staff. No confirmations or endorsements are intended or implied.

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

STRUCTURAL COATINGS IN ALUMINUM ALLOY MICROTRUSS MATERIALS

Bosco Yu, Glenn D. Hibbard

Department of Materials Science and Engineering, University of Toronto, 184 College Street, Toronto, Ont., M5S 3E4, Canada

Keywords: Surface Engineering, Aluminum, Microtrusses, Energy Absorption

Abstract

Incorporating an internal cellular architecture of open space is one strategy to increase the potential functionality of aluminum alloys. Stretch-dominated microtruss cellular architectures, which are designed such that externally applied loads are resolved axially along the internal struts, provide enhanced strength and stiffness at low densities when compared to conventional metal foams. In this study we introduce the idea of using a structural coating to reinforce AA3003 aluminum alloy microtrusses. Because the internal surface area is large and the strut cross-sectional dimensions can be as small as hundreds of microns, only a 40 urn thick hard anodized alumina coating was needed to induce a four-fold increase in compressive strength and a six-fold increase in energy absorption at virtually no additional weight penalty. Backscatter electron microscopy was used to examine the failure mechanisms of the structural coatings and the cores in order to explain such a large change in behaviour.

Introduction

Microtrusses cellular materials have attracted considerable interest because of their structural efficiency and their potential to exhibit multifunctional characteristics [1-10]. In term of load bearing capacity, microtrusses are designed with a high degree of internal strut connectivity. Externally applied loads are resolved by axial (stretching) rather than bending deformation, making them up to ten times stiffer and three times stronger than conventional metal foams at low relativity densities [11], Since the internal surface areas are large and the strut cross-sections are small, surface modification can also be used to further enhance their mechanical behaviour. Recently, electrodeposition of a ~50um thick coating of ultrahigh strength nanocrystalline nickel was found to double and triple the compressive strength of low carbon steel and aluminum microtrusses, respectively [5, 6, 12].

Inelastic buckling was the overall failure mechanism seen for both conventional and structurally coated microtrusses [6]. Under this failure mechanism, the load carrying capacity of the microtruss progressively decreases making these materials less desirable for energy absorption applications. A recently study, however, indicated the possibility of enhancing the energy absorption behaviour of aluminum microtrusses by applying anodic oxide coating in order to change their overall collapse mechanism from inelastic buckling to hinge rotation and fracture [13]. The present study examines this hinge rotation and fracture failure mechanism in order to understand the mechanism of enhanced energy absorption in anodized microtruss materials.

3

Experimental

AA3003 aluminum perforated sheets (sheet thickness of 0.74 ± 0.01 mm) were purchased from McNichols Perforated Products (Altanta, GA). The perforation pattern had square holes of 5.10x5.10 mm punched on a square lattice of 6.36x6.36 mm creating an open area fraction of <P = 0.64; Fig. la shows a schematic diagram of the starting sheet materials. The aluminum sheets were annealed at 600 °C for 30 minutes and then water cooled before forming into pyramidal microtrusses, after Bouwhuis et al. [5, 14]. Three forming/annealing cycles (with intermediate annealing treatments of 600 °C for 30 minutes) were used to achieve an internal truss angle of 0 = 42.6° ± 0.1° and a relative density of p = 4.8% (in Fig. lb). The microtrusses were subsequently given a post-fabrication annealing treatment of 600 °C for 30 minutes to remove any fabrication-induced work hardening, which was confirmed by taking microhardness measurements on axial cross-sections of the microtruss struts (an average of 36 ± 2 HV over all sample types using 0.98 N applied load and 10 s dwell time).

Fig. 1. Schematic diagram of a starting perforated aluminum sheet (a) and the final architecture after forming (b).

The AA3003 pyramidal microtrusses were then subjected to three surface cleaning processes (degreasing, etching, and de-smutting) followed by hard anodizing to form a 40 urn thick anodic aluminum oxide coating on the surface of the microtruss cores. Samples were rinsed with de-ionized water between each step and after hard anodizing. Table I summarizes the process conditions for each step. The microhardness of the anodized coating was 537 ± 38 HV (using 0.98 N applied load and 10 s dwell time). Coated and uncoated (only subjected to surface cleaning) samples were then tested in uniaxial compression at a cross-head displacement rate of 1 mm/min using confinement plates that restricted the nodal displacement in the microtrusses. In addition, four coated samples were partially compressed to different strain level (starting from e = 0 to e = 0.53). The axial cross-sections of their internal struts were characterized by scanning electron microscopy (SEM).

4

Process Name Degreasing

Etching

De-smutting

Hard anodizing

Table I. Surface Cleaning and Hard Anodizing Specificatior Solution

5g/l Na2C03

5g/lNa3P04

5wt% NaOH

20vol% HNOj

12wt% H2SO4 lwt% (COOH)2

Time 10 min

2 min

1 min

45 min

Temperature 87 ± 5°C

57 ± 2°C

33 ± 2°C

10±1°C

Current density

4 A/dm2

Results and Discussion

Figure 2 presents typical compressive stress-strain curves of the coated and uncoated samples. Both samples exhibit an initial elastic region (of slope E) before reaching a peak stress (o>ra*) that is followed by post-peak softening until a valley stress (o\,a//«y), and eventually final densification is reached.

Fig. 2. Typical compressive stress-strain curves of the coated and uncoated microtrusses.

Figure 3 plots the slope of the stress-strain curve as a function of strain for both sample types. At a low strain level, both the coated and uncoated microtrusses elastically deform with nearly constant compressive modulus. After a strain of about e = 0.05, the samples begin to deform plastically and the moduli values decrease, eventually dropping below zero. In the case of the uncoated samples, the modulus drops gradually to zero; the first ̂ -intercept is the strain when the peak strength is recorded (e^*) while the second ^-intercept is the strain when the valley strength is recorded (£«,//*,,). In the case of the coated samples, the local modulus passes through large oscillations as it progressively decreases; the firsts-intercept is lower than the value of the Epeak while the last ^-intercept is higher than the value of the emuey- These modulus fluctuations indicate that the coated microtrusses are subjected to series of fracture events as they collapse, which is not commonly seen in the inelastic buckling of conventional microtrusses.

5

Fig. 3. Modulus (do/de) curves of the coated and uncoated microtrusses.

The energy absorption properties of the microtrusses can also be examined by measuring the area under their stress-strain curves (starting from half the peak stress to twice the peak stress, after [15]):

çc(Zaveak)

Jdensification = 1 n ■. ( '

The compressive properties of the coated and uncoated microtrusses are summarized in Table II. It seems at first surprising that with only a 40 um thick anodized coating, the coated microtrusses were twice as stiff, four-times as strong, and had six-times the energy absorption capability of the uncoated microtrusses. Note that the post-peak softening region (defined as the difference between epeak and evaney) of the coated samples is 57% higher than that of the uncoated samples. Prolonging the post-peak behaviour while maintaining a high stress level is the key reason why the coated microtrusses have superior energy absorption characteristic.

Table II. Compressive Properties of the Coated and Uncoated Microtrusses. Sample

Coated

Uncoated

E [MPal

33

15

Gpeak

[MPal 4.20

1.03

Speak

[mm/mm] 0.26

0.1

Gvalley

[MPal 3.19

0.58

^valley

[mm/mm] 0.62

0.33

Jdensification

[MPal 2.96

0.50

Macroscopic Failure Mechanism

While the uncoated AA3003 microtrusses follow a conventional inelastic buckling collapse mechanism, the coated microtrusses fail in a very different manner. Backscatter scanning electron microscopy (figure 4) shows the evolution of the coated axial cross-sections of the coated microtrusses (^-direction in figure lb) after different levels of compressive strain: e = 0 (a), e = 0.24 (b), e = 0.33 (c), e = 0.44 (d), and e = 0.53 (e). The left-hand side of the as-coated sample (figure 4a) shows that the top part of the node is flattened due to the fabrication pin-indentation while the bottom part has a convex curvature, creating a region with a smaller cross-sectional thickness between the node and the strut. At e = 0.24 (figure 4b), the microtruss is compressed closed to its peak strength. The bottom part of the node is now flattened due to compression and the top part is bulged outward. On the right-hand side closer to the node, a series of oxide pulverizations is seen, which explains the modulus fluctuation of the coated sample (figure 3). At strain-levels beyond the peak strain (e = 0.33, 0.44 and 0.53), pulverization continues from the edge of the node, expanding along the strut length. Figure 5 shows the axial cross-sectional thickness of the core for each of these strain-levels along the length of the strut. For strain values greater than the initial peak strain of e = 0.26, the samples show that the axial cross-section thickness first decreases to a minimum at the edge of the node, then increases within the pulverized region before decreasing again to the thickness level seen in the undamaged region of the strut. It is interesting to note that the strut remained at an angle of-40° (near the starting angle of 8 = 42)6 throughout the collapse sequence. The fact that the undamaged segments of the struts retain nearly the same angle throughout collapse is one of the contributing factors to the prolonged post-peak behaviour of the coated microtrusses.

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Fig. 4. Backscatter scanning electron micrographs of the coated microtrusses at compressive strains of e = 0 (a), e = 0.24 (b), e = 0.33 (c), e = 0.44 (d), and e = 0.53 (e).

Fig. 5. Axial cross-sectional thicknesses of coated microtrusses cores at different strain levels starting from the centre of the node at increment of lOOum.

8

Microscopic Failure Mechanism

Figure 6 compares the microscale differences between uncrushed (a) and pulverized (b) microtrusses. The manganese-rich second phase in the AA3003 alloy is shown as the elongated white particles in the micrographs. Before compression, the second phase particles are aligned parallel to the forming direction of the strut. As pulverization progresses, the anodized coatings penetrate into the aluminum core, forcing the aluminum to undergo extensive plastic deformation and causing the second phase particles to follow a wave-like flow pattern (traced by white arrows). It is this combination of coating fracture and extensive core plastic deformation which allows the stress to be maintained at the relatively high levels seen during post-peak deformation, resulting in the observation of enhanced energy absorption.

Fig. 6. High magnification backscatter scanning electron micrographs of the uncrushed (a) and pulverized (b) microtrusses.

Conclusions

Structural coated ceramic/metal microtruss composites were made by hard anodizing AA3003 aluminum microtruss cores. With only a 40 urn thick hard anodized alumina coating, the coated microtrusses were twice as stiff, four-times as strong, and had six-times the energy absorption capability of the uncoated microtrusses. At the microscopic level, hard anodized coatings continuously penetrate into the softer aluminum cores, forcing the aluminum to undergo significant amounts of plastic deformation. At a macro-scale, coating pulverization allows the node to expand without buckling. As a result, the post-peak behaviour of the coated microtrusses is prolonged while the stress is maintained at a high level. Consequently, the ceramic/metal microtruss composites are able to exhibit superior energy absorption characteristics compared to conventional microtrusses. Studying the unconventional failure mechanisms of these ceramic/metal microtruss composites will allow strut design and optimization of these new light-weight composites for both structural and energy absorption applications.

9

References

[I] V.S. Deshpande, N.A. Fleck, and MF. Ashby, "Effective properties of the octet-truss lattice material," Journal of the Mechanics and Physics of Solids, 49 (8) (2001), 1747-1769. [2] S. Chiras et al., "The structural performance of near-optimized truss core panels. International Journal of Solids and Structures, 39 (2002), 4093-4115. [3] N. Wicks and J.W. Hutchinson, "Performance of sandwich plates with truss cores," Mechanics of Materials, 36 (2004), 739-751. [4] A.J. Jacobsen , W. Barvosa-Carter, and S. Nutt, "Compression behavior of micro-scale truss structures formed from self-propagating polymer waveguides," Acta Materialia, 55 (20) (2007), 6724-6733. [5] B.A. Bouwhuis et al., "Structural nanocrystalline Ni coatings on periodic cellular steel," Composite Science and Technology, 69 (3-4) (2009), 385-390. [6] E. Bele, B.A. Bouwhuis, and G.D. Hibbard, "Failure mechanisms in metal/metal hybrid nanocrystalline microtruss materials," Acta Materialia, 57 (19) (2009), 5927-5935. [7] D.D. Radford, N.A. Fleck, and V.S. Deshpande, "The response of clamped sandwich beams subjected to shock loading," International Journal of Impact Engineering, 32 (6) (2006), 968-987. [8] K.P. Dharmasena et al., "Dynamic compression of metallic sandwich structures during planar impulsive loading in water," European Journal of Mechanics - A/Solids, 29 (2010), 56-67. [9] C.J. Yungwirth et al., "Explorations of Hybrid Sandwich Panel Concepts for Projectile Impact Mitigation," Journal of the American Ceramic Society, 94 (SI) (2011), S62-S75. [10] K.P. Dharmasena et al., "Response of metallic pyramidal lattice core sandwich panels to high intensity impulsive loading in air," International Journal of Impact Engineering, 38 (5) (2011), 275-289. [II] V.S. Deshpande, M.F. Ashby MF, and N.A. Fleck, "Foam topology bending versus stretching dominated architectures. Acta Materialia, 49 (2001), 1035-1040. [12] K. Abu Samk, B. Yu, and G.D. Hibbard, "Architectural design in stretch-formed nanocrystalline microtruss composites" (Manuscript number: JCOMA-11-890, Composite Part A: Applied Science and Manufacturing, 2011). [13] Bele E et al., "Structural ceramic coatings in composite microtruss cellular materials," Acta Materialia, 59 (15) (2011), 6145-6154. [14] B. Bouwhuis and G.D. Hibbard, "Failure Mechanisms during Periodic Cellular Metal Fabrication by Perforation Stretching," Metallurgical and Materials Transactions A, 39 (12) (2008), 3027-3033. [15] B. Bouwhuis and G.D. Hibbard, "Compression testing of periodic cellular sandwich cores," Metallurgical and Materials Transactions B, 37 (6) (2006), 919-927.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

LASER CLADDING OF HIGH-PERFORMANCE CPM TOOL STEELS ON HARDENED H13 HOT-WORK TOOL STEEL FOR AUTOMOTIVE

TOOLING APPLICATIONS

Jianyin Chen, Lijue Xue

Centre for Automotive Materials & Manufacturing Industrial Materials Institute, National Research Council of Canada

800 Collip Circle, London, Ontario, Canada N6G 4X8

Keywords: Laser cladding, CPM tool steels, H13 tool steel, Tooling

Abstract

This paper summarizes our work on laser cladding of high-vanadium carbide CPM tool steels (3V, 9V and 15V) onto hardened chromium hot-work AISI HI3 tool steel with a hardness of HRc 50-55 in order to substantially enhance abrasive wear resistance, which provides a great potential for fabricating high-performance automotive tooling at affordable cost.

Introduction

For many years, automotive manufacturers are subjected to significant challenges to reduce the cost of auto parts to maintain their global competitiveness, in which tooling constitutes a substantial portion of production cost (such as plastic injection/blowing molding, die casting, forging, hydroforming, stamping, and others). Laser cladding can deposit a high-performance material onto a cheap base material (or substrate) to produce a metallurgical^ sound and dense clad [1, 2]. This process can be used to enhance the performance of automotive tooling fabricated by a low-cost substrate, in order to extend tooling life, reduce the machine down-time, and save overall tooling cost (consequently the per-part cost), and has attracted lots of interests. In addition, laser cladding uses a relatively low heat input to minimize distortion of the tools being cladded, produces a functional clad with a metallurgical bonding to the base material (or substrate), and minimizes the dilution of a cladding material into a substrate which, subsequently, reduces undesired deterioration of properties of the base material (or substrate).

The objective of this research aims at the development of laser cladding process to deposit highly wear-resistant Crucible's CPM 3V, 9V and 15V tool steels onto a low-cost hardened AISI H13 base material (or substrate). H13, a kind of chromium hot-work tool steel, is widely used to make automotive tooling. Even though the hardened H13 steel with a hardness of HRc 50~55 could provide a reasonably good wear resistance, excessive wear is still inevitable after a period of operation. It is highly desired that wear-resistant materials can be cladded at the "easy-to-wear" location(s) of the HI 3 made tools to enhance their mechanical performance. CPM tool steels, with an excellent combination of toughness, hardness and wear resistance [3], could provide such an opportunity to overcome current deficiencies of H13 tool steel. However, CPM steels (with a hardness of more than HRc 58) have posted technical challenges to laser cladding, which could easily crack the clads and/or substrates during cladding. So far, a number of works [4-8] on the subject have been published. This paper summarizes our investigation of laser cladding of

11

3V, 9V [7] and 15V [6] tool steels on hardened H13 base material. The microstructural evolution and mechanical performance of the clads so obtained were discussed as well.

Experimental

Cladding and Base Materials

The gas-atomized spherical CPM powders (Crucible Research Center, Pittsburgh, PA) with a size of 10-45 urn were used as cladding material (Table I). About 17 mm thick hardened wrought H13 steel plates with a hardness of HRc 50~55 were used as base material (or substrate).

Table I. Chemical compositions of CPM powders and wrought H13 substrate (wt. %)

Element 3V 9V 15V

H13 [91

Cr 7.49 5.24 5.23 5.10

V 2.59 9.10 14.44 0.61

Mo 1.35 1.27 1.39 1.11

W 0.14 0.01

--

Si 0.89 0.89 0.92 1.00

Mn 0.39 0.49 0.41 0.47

C 0.83 1.82 3.5 0.39

Ni

---

0.11

Cu

---

0.09

P 0.010 0.011 0.022 0.02

Fe Bal. Bal. Bal. Bal.

Laser Cladding Process

Laser cladding process by blown powder feeding was selected for the current study (Fig. la). In the setup, a 500 W pulsed Nd:YAG laser, a precision powder feeder and a computer numerically controlled (CNC) motion system were used. The laser focusing optics and powder feeding nozzle were mounted on the Z-axis of the CNC motion system while a hardened H13 plate as a substrate was clamped on the X-Y motion table. Laser cladding was carried out in a glove box filled with argon to maintain an oxygen level below 50 ppm, and in addition, the argon was also used as a carrying gas for the delivery of CPM powder and a shielding gas for the laser optics during the process.

(a) (b) Figure 1. (a) Schematic diagram of laser cladding process by blown powder feeding, and (b)

metallurgical sound laser-clad 9V coupons.

During cladding, a laser beam with an average power of 100-300 W was defocused onto the substrate to create a melt pool (about 1.0 mm in diameter) along with the injected CPM powder. With a traverse speed of 5~15 mm/s, a track of the molten cladding material was deposited onto the substrate and rapidly solidified to form a single clad track that was metallurgically bonded to

12

the substrate. By re-melting a portion of the preceding clad track (with an overlap of 25-35%) and the substrate, along with injected powder, the second clad track, which neighbored to the first one, was deposited. The 1st layer of the clad was formed by depositing multi clad tracks with a partial overlap. Subsequent layers can be further deposited onto the previous layers whereby they were effectively the new substrate, until the desired clad thickness (hereby, about 0.6-0.8 mm) was achieved (Fig. lb).

In general, for laser cladding of highly alloyed tool steels, cracks often occur in the clad as a result of the build-up of residual tensile stresses mainly induced by thermal mismatch between the clad and the cold substrate. As one of many possible solutions to address this issue, prior to any cladding, the substrate can be preheated to certain temperature in order to alleviate its thermal mismatch with the clad and control martensitic transformation in the clad when the clad cools after it re-solidifies. This might subsequently lessen the tensions in the clad [10]. Consequently, the cladability of the cladding material is improved. Concerning the current case, the hardened H13 substrate has been heated to a temperature below its maximum tempering temperature (540-550 °C) during cladding. Moreover, laser process parameters have also been optimized to obtain a fully dense clad. Afterwards, a double tempering was applied on the clad specimens at a temperature range of 150-600 °C to progressively eliminate retained austenite in the clads and promote further precipitation of fine carbides.

Materials Characterization

An Olympus optical microscope (OM) and a Hitachi scanning electron microscope (SEM) were used for the current metallographic study. A Philips X'Pert X-ray diffractometer (XRD) with graphite-monochromatic CuKa radiation was used to identify phases in the clads. The hardness of the clads were measured using a Rockwell hardness testing system, and the microhardness profiles of the laser-clad CPM steels were measured on a Buehler's microhardness tester using a Vickers indentor.

Results and Discussion

Metallurgical Soundness

All laser-clad 3V, 9V and 15V tool steels show an acceptable metallurgical soundness even though a few tiny gas-trapped pores might be observed in the clads. Normally, the depths of penetration (or dilution) of the CPM tool steels into the H13 substrate were controlled to be below 0.15 mm, and the thicknesses of the heat-affected zones (HAZs) were below 0.5 mm.

Phases

Figure 2 is the X-ray diffractograms of "as-clad" and "double tempered" 3V, 9V [7] and 15V [6] clads. The primary carbides in the 9V and 15V clads were mainly vanadium carbides (VC), while the presence of carbides (vanadium and chromium carbides) in the 3V clad cannot be detected due to a low XRD resolution. In addition to various carbides, the major phases in all three clads were martensite (a) and retained austenite (y1). Comparatively, the "as-clad" 9V steel possessed less y' phase than the "as-clad" 3V and 15V steels did (Fig. 3).

13

Figure 2. X-ray diffractograms of laser-clad (a) 3V, (b) 9V and (c) 15V tool steels in "as-clad" and "double tempered" conditions.

The y' phase in the clads is usually considered as a thermally unstable phase and may be transformed into a phase during their service, causing dimensional change, brittleness or even cracking. Moreover, tensile stresses could present in the clads since thermal shrinkage in the solidified clads during cooling is constrained by the unaffected substrate. Although any martensitic transformation in the clads during cooling, which induces a volumetric expansion, could partially or fully cancel the tension [11], the presence of an excessive amount of y1 phase in the clads, nevertheless, adversely dampens such kind efforts. Thus, the role of a post tempering on the CPM clads is not only to eliminate y1 phase but also to reduce tensile stresses as well [10].

Figure 3 summarizes the variations of relative quantities of y' phase, which were characterized by I/(3ii/Ia(2ii), and the hardness in the 3V, 9V and 15V clads with double tempering temperature. With a rise in temperature up to 500-600 °C, the amounts of y1 phase in the clads gradually disappeared. It is interesting to note that double tempering before 400 °C had no meaningful effect on the overall hardness of the CPM clads. Comparatively, the hardness of the CPM clads could achieve their peak values ("secondary hardening") after double tempered at 500~550 °C.

Figure 3. Variations of retained austenite (y') and hardness of laser-clad 3V, 9V and 15V tool steels with double tempering temperatures.

Macro- and Microstructures

Figure 4 exhibits SEM morphologies in "as-clad" 3V, 9V and 15V solidified melts cross-sectioned transverse to the laser cladding direction: columnar dendritic starting from the border of two neighbored melts, growing and transiting into equiaxed cellular near the top of the melt (laser-clad 3V), or dominant equiaxed cellular within the melts (laser-clad 9V and 15V).

14

Figure 4. Solidification morphologies of "as-clad" (a) 3V, (b) 9V and (c) 15V solidified melts.

To explain the observed columnar dendritic to equiaxed cellular in a solidified melt, Hemmati et al. [12] stated that solid-liquid interface growth rate R (mainly a direct function of cladding speed) and temperature gradient G in the melt control the solidification structure, where the G/R ratio determines the solidification mode and the GR product controls scale of the solidified structure; by decreasing the G/R ratio, solidification structure could change from planar to cellular and then columnar dendritic and equiaxed cellular. In multi layer deposition, heat accumulation (mainly controlled by laser heat input and latent heat released by solidification), which gradually increases the size of a melt during deposition, reduces the G typically associated with columnar dendritic in the melt and provides an opportunity for a transition from columnar dendritic to equiaxed cellular once nucleation of equiaxed cellular grains occur in the liquid ahead of the columnar dendrites [12]. The solidified morphology in the 3V melts could be a result of the similar reasons. For the laser-clad 9V and 15V tool steels, a possible relatively low G (caused by relatively high heat input and long pulse duration) could already have induced dominated equiaxed cellular in their solidified melts.

Furthermore, laser cladding could present a heterogeneous microstructure that differs from point to point, since cladding is performed through a partial overlap of clad tracks to cover a desired area on the surface of a substrate. When a new clad track is deposited, a portion of previously deposited neighbored clad track is re-melted and solidified together with the injected powder to form the new clad track; while the microstructure in the previously deposited clad track adjacent to the re-melting area is severely thermally affected. For the steels, these local thermal cycles could activate solid-state phase transformations that lead to progressive modification of microstructures of the previously deposited clad tracks such as martensite tempering, partial or total austenitization as well as transformation of austenite into martensite [13-15]. Therefore, two distinct but adjacent zones could alternatively exist in "as-clad" material: "as-deposited clad zone (ADCZ)" and "re-heated zone (RHZ)" [5]. The "RHZ" in laser-clad steels, more accurately, could be divided into two sub-zones (or regions): (1) a "re-austenitized region (RAR)" in which, as the re-heating temperature exceeds Aci, martensitic phase in that region would transform into austenite (which would transform into martensite again upon subsequent rapid cooling); and (2) a "tempered region (TR)" where since the re-heating could only reach to the temperature lower than Ad, certain level of tempering in that region might occur [13-15]. Figure 5a schematically depicts the morphological relationship of "ADCZ", "RAR" and "TR" [10]. For the laser-clad 9V and 15V tool steels, the morphological difference of the "ADCZ" and the "RAR" can be observed while it was hard to differentiate the "ADCZ" with the "TR" for the 3V, 9V and 15V tool steels since highly alloyed CPM tool steels have a high resistance against tempering. Hence, under the circumstance of rapid heating and subsequent cooling, self-tempering induced by those

15

thermal cycles during the period of cladding may not sufficiently lead to observable modification of microstructures in the "TR".

Figure 5. (a) Schematic diagram of multi overlapped clad beads (tracks) cross-sectioned transverse to the laser cladding direction, showing "as-deposited clad zone (ADCZ)", "re-

austenitized region (RAR)" and "tempered region (TR)"; SEM microstructures near the border of two neighbored (b) 3V and (c) 9V solidified melts, respectively; (d) in "ADCZ" of a 15V

solidified melt, and (e) near the border of two neighbored 15V solidified melts.

In terms of laser-clad 3V tool steel, the "as-clad" 3V microstructure along the border of two neighbored solidified melts reveals that re-austenitization in the "RAR" around cellular area caused an insignificant modification of the morphology (Fig. 5b). Moreover, it can be estimated from the same figure that the primary dendritic arm spacing (PDAS) and the secondary dendritic arm spacing (SDAS) in "ADCZ" of a solidified melt were about 2-3 and 1-2 urn respectively; while the cellular size (CS) in the solidified melt was about 2-4 um. Similarly, the dominant cellular structure in a 9V solidified melt (Fig. 5c) reveals that re-austenitization in the "RAR" caused certain modification of the cellular microstructure: the eutectic carbides at interdendrites gradually developed into coarse globular carbide particles. It can be estimated that the CS in a 9V solidified melt was about 2~4 ^m. It is no doubt from the scale of the solidified features that both laser-clad 3V and 9V tool steels experienced a rapid solidification. Figure 5d exhibits SEM microstructure of "as-clad" 15V solidified melt in "ADCZ": multi-shaped primary vanadium carbides (globular and star-like) were dispersed on some equiaxed cellular matrix (with a CS of about 4-7 um) as a form of bands guided by stirred molten flows; while some cellular matrix shows much less presence of primary carbides. Kury et al. [16], having investigated rapidly solidified gas-atomised powder of highly alloyed V-Cr tool steel (3%C-3%Cr-12%V), noted that the morphological variants of carbides in tool steels of ledeburite type are closely related to complicated thermal condition during their solidification. A further study is needed to explain the current morphologies of carbides in the laser-clad 15V solidified melts. Comparatively, re-austenitization in the "RAR" (Fig. 5e) has caused some noticeable change: partial disappearance

16

of cellular characteristics in the matrix and precipitation of more fine globular carbides from the matrix along with existing primary carbides.

Microhardness Profiles

The microhardness profiles of laser-clad 3V, 9V [7] and 15V [6] tool steels, in both "as-clad" and "double tempered" conditions, are depicted in Fig. 6. In "as-clad", the average microhardness of the hardened H13 substrate was about 509 HV while that of the 3V, 9V and 15V clads were about 601 HV, 646 HV and 740 HV, respectively. No significant "soft" zone was observed in the HAZ just beneath the 9V and 15V clads, although a slight drop of hardness in the HAZ was observed in the 3V clad. The average microhardness of the 3V and 15V clads after double tempered at 540-550 °C were substantially increased to about 682 HV and 833 HV, respectively; while that of the 9V clad after double tempered at 500 °C was increased to 670 HV as well. Meanwhile, the hardness of the HAZ in the 3V clad has been restored to the same level as the unaffected H13 substrate. The increases in the hardness of CPM clads (or "secondary hardening") after double tempering in the range of 500~550 °C, was ascribed to the precipitation of fine carbides on the matrices, which is frequently observed in highly alloyed tool steels after a proper tempering treatment.

Figure 6. Microhardness profiles of laser-clad (a) 3V, (b) 9V [7] and (c) 15V [6] tool steels in both "as-clad" and "tempered" conditions (HAZ: heat-affected zone; D: the penetration zone).

Furthermore, it was found in Fig. 3 that performing a double tempering below certain temperature, the CPM clads retained almost the same hardness, whereas with a double tempering around 500-550 °C, the hardness of the clads reached their "secondary hardening" with a maximum value of about HRc 62, 65 and 69 for the 3V, 9V [7] and 15V [6] clads, respectively, which were either equivalent to or higher than their corresponding P/M CPM tool steels [3]. This phenomenon might be ascribed to process-induced refined microstructures in the clads. Comparatively, the hardened H13 substrate could retain a hardness of about HRc 53 until after a double tempering beyond 540-550 °C. Therefore, the heat treatments selected for the laser-clad CPM steels, if less than 540-550 °C, would not compromise mechanical performance of the hardened H13 base material. Furthermore, after the "secondary hardening", the hardness of the laser-clad CPM tool steels would gradually drop but were still higher than the corresponding P/M CPM tool steels [3]. Nevertheless, a proper selection of tempering temperature could lead to a different hardness in laser-clad CPM tool steels which, in return, could meet different mechanical performance needs in various tooling applications such as wear resistance, machinability and toughness [3].

17

Summary

Laser cladding with blown powder feeding has successfully produced metallurgically sound CPM clads on hardened H13 base material (or substrate), which could be potentially applied for automotive tooling manufacturing and repair. The metallurgical characteristics of the investigated clads can be summarized as follows:

(1) "As-clad" 3V, 9V and 15V tool steels show fine columnar dendritic and/or equiaxed cellular structures containing considerable amount of martensite, retained austenite and carbides. The hardness of "as-clad" 3 V, 9V and 15V tool steels can reach HRc 58, 65 and 65, respectively.

(2) After double tempered at 500-550 °C, the retained austenite in the clads completely disappears; while the hardness of laser-clad 3V, 9V and 15V tool steels reach their "secondary hardening", with a hardness of HRc 62, 65 and 69 respectively, which are either equivalent to or higher than the corresponding P/M-made CPM steels; whereas the hardness of H13 substrate with a hardness of about HRc 50-55 still retains almost the same even after double tempered up to 540-550 °C. Therefore, the post-cladding tempering selected for the current laser-clad CPM tool steels would not compromise the mechanical performance of the hardened H13 substrate.

Acknowledgments

Authors would like to thank Mr. Jon Fenner, Technical Officer, for his preparation of laser-clad CPM 3V, 9V and 15V tool steel coupons for the current metallurgical investigation.

References

[I] W.M. Steen, Laser Materials Processing (Td Ed.) (London: Springer, 1998), 248-251. [2] R. Vilar, J. Laser Appi, ll(2)(1999),64-79. [3] Crucible Service Centers, Data Sheets for Crucible CPM® 3V®, 9V®, & 15V® tool steels [4] Y.P. Hu, C.W. Chen, K. Mukherjee, J. Mater. Sei., 33(5)(1998),1287-1292. [5] S.-H. Wang, J. Chen, L. Xue, Surf. Coat. Techno!., 200(11)(2006), 3446-3458. [6] J. Chen, L. Xue, H. Visscher, J. Wolfe, in Proc. Surface Protection for Enhanced Materials Performance, MS&T09, Oct. 25-29, 2009, Pittsburgh, PA, pp.2090-2101. [7] J. Chen, L. Xue, in Proc. High Performance Tooling Materials, MS&T'IO, Oct. 17-21, 2010, Houston, TX, pp.2459-2470. [8] J. Leunda et al., Physics Procedia, 12(2011), 345-352. [9] ASM International, TLS™ H13 PQ (Filing Code: TS-588 Tool Steel), Alloy Digest, June 2001. [10] J. Chen, S.-H. Wang, L. Xue, J. Mater. Sei., (2011), DOI 10.1007/sl0853-011-5854-4. [II] M. Pilloz, J.M. Pelletier, A.B. Vannes, J. Mater. Sei., 27(1992), 1240-1244. [12] I. Hemmati, V. Ocelik, J. Th. M. De Hosson, J. Mater. Sei., 46(2011), 3405-3414. [13]L. Costa et al., Mater Sei Forum, 473^174(2005), 315-320. [14] L. Costa et al., In: D. Keicher, J.W. Sears, J.E. Smugeresky (eds), Proc. Int'l Conf. on Metal Powder Deposition for Rapid Manufacturing, Princeton (NJ): MPIF; 2002, pp. 172-179. [15] R. Vilar, R. Colaco, A. Almeida, Opt. Quant. Electron, 27(1995), 1273-1289. [16] M. Kusy et al., Mater. Sei. Eng. A, 375-377(2004), 599-603.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ELECTRON BEAM DEPOSITED MULTILAYER OPTICAL INTERFERENCE COATINGS USING OXIDE COMPOSITES

Ankush N Nayak1, N K Sahoo2, R B Tokas2, Arup Biswas2, Nitin M Kamble2

'Department of Metallurgical and Materials Engineering, National Institute of Technology Kamataka, Surathkal;

P. O. Srinivasnagar; Mangalore, Karnataka, 575025, India

2Applied Spectroscopy Division, Bhabha Atomic Research Centre; Trombay; Mumbai, Maharashtra, 400085, India

Keywords: optical coatings, thin films, electron beam physical vapor deposition, high-reflection (HR) mirror, zirconium dioxide, magnesium oxide, aluminium oxide, atomic force

microscopy, high power laser mirror.

Abstract

Optical multilayer interference coatings are not only the key elements/components of the lasers, synchrotron (beam lines), and solar devices but also serve to propagate, deliver and manipulate electromagnetic radiations involved there for materials science experiments. Composite oxide thin film materials have added several promising dimensions with respect to the design, development of such precision devices related to such applications. Binary Zr02-MgO and ternary Zr02-MgO-Al2C>3 oxide composite thin films have been deposited using electron beam physical vapor deposition (EB-PVD) technique and nano-metric multilayer devices utilizing such films in a regular periodic layered design have been developed. As a specific objective, a multilayer high-reflection (HR) laser mirror having a designated bandwidth has been designed and developed for the Nd:YAG second harmonic laser wavelength of 532 nm. These composite thin films and multilayers have been characterized using various microstructural probing techniques.

Introduction

Optical multilayer interference coatings have potential applications in various technologies and science researches involving various electromagnetic sources, viz., lasers, synchrotrons, neutron generators and laboratory spectral sources. Thin films of dielectric materials, especially refractive oxides have been of great research interest owing to their wide range of optical properties and specifications. The materials used in this study are oxide composites of binary ZrÛ2-MgO and ternary Zr02-MgO-Al203 compositions. These materials have been subject to various research investigations and their thin films have been deposited individually and studied [1, 2]. However, there is a lack of a complete knowledge available in the literature concerning their microstructural and optical properties and their interdependences especially in thin film phases. In the present study, several single layer thin films of the binary and ternary composites have been deposited using reactive electron beam deposition process. These single layers of five quarter-wave optical thickness at monitoring X of 550 nm (-600 nm, physical thickness) have been deposited for process optimization studies under various oxygen partial pressures and a fixed deposition rate of 5 À/s. The thin films have been characterized for the refractive index, absorption coefficients, and physical thicknesses by UV-VIS-NIR Spectrophotometry. Subsequently, utilizing optimized thin film composites in a periodic design, a multilayer high-reflection (HR) laser mirror having a specific desired bandwidth of -60 nm has been designed and developed for the Nd:YAG second harmonic laser wavelength of 532 nm. This

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multilayer mirror consists of 51 alternate layers of binary and ternary composite oxide thin films to achieve the desired design specifications. All these composite thin films and multilayers have been characterized using various techniques like Spectrophotometry, Atomic Force Microscopy (AFM), and Grazing Incidence X-Ray Reflectivity (GIXR).

Conceptual Basis of Optical Interference Multilayers

A basic type of optical thin film multilayer geometry is a stack of alternate high- and low-index films, all one-quarter-wavelength thick. Light reflected within the high-index layers does not undergo any phase change on reflection, whereas that reflected within the low-index layers undergoes a change of phase of 180°. It is quite easy to see that the various components of the incident light produced by reflection at successive boundaries throughout the assembly will reappear at the front surface all in phase so that they will recombine and interfere constructively. That's why such a multilayer geometry is termed as interference optical coatings and this means that the effective reflectance of the assembly can be made very high just by increasing the number of layers by enhancing the inference components. This is the basic form of the high-reflectance coating. When such a coating is constructed, it is found that the reflectance remains high over only a limited range of wavelengths, depending on the ratio of high and low refractive indices. Outside this zone, the reflectance changes abruptly to a low value. Because of this behavior, the quarter-wave stack is used as a basic building block for many types of thin-film filters. An essential common feature of dielectric optical materials is their very low absorption (i.e., absorption coefficient a< 103cm"') in some relevant portion of the spectrum; in this region they are essentially transparent [3]. The refractive index n is basically the only optical constant of interest so far as optical coating multilayer design is concerned. Since the high reflectance of a single AJ4 film is due to the constructive interference of the beams reflected at both surfaces, the effect can be enhanced by phase agreement in the reflected beams from multiple film layers. What is required is a stack of alternating high (H) and low (L) index A/4 films. Next to the substrate is the usual high index layer so that the stacking order is HLHLHLHL...up to z"1 layer. For z layers it has been calculated that the maximum reflectance is given by [4]:

where w#, ni, and r\2 are the high, low, and substrate indices respectively. The magnitude of the reflectance increases with the number of layers. The number of sideband oscillations outside the high-reflectance zone also increases with number of layers. The spectral width of the high reflectance zone is a function of the ratio of the refractive indices of the involved films. Thus, by reducing the difference between KH and n i ; a narrow-band reflection filter can be generated. Multilayer dielectric interference systems are ideally suited as reflection coatings for fully reflecting and partially transmitting laser mirrors. Negligible absorption means that reflectance of almost 100% can be achieved.

Experimental

Single-layer thin films of the binary ZrO^-MgO and ternary Zr02-MgO-Al203 composites were deposited, using reactive electron beam physical vapor deposition technique, on a quartz substrate at various oxygen partial pressures and at substrate temperatures of 150°C and 300°C. Each of these films had an optical thickness of five to six quarter-wavelengths. The

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rate of deposition was maintained at 5 À/s throughout the deposition process. The development of the high-reflection laser mirror involved deposition of 51 alternate layers of high- and low-index composites, i.e., Zr02-MgO and Zr02-MgO-Al203 respectively, to meet design specifications. The oxygen partial pressure was maintained at 8xl0"3 mbar and the substrate temperature was 300°C. The deposition rate was 5 A/s and each composite layer was one quarter-wavelength thick. Electron beam deposition is well suited for synthesis of thin films of refractory oxides mainly because of the high temperatures attainable. The impinging electrons heat up only the surface and the material in contact with the crucible remains solid. Effectively, the molten material is contained in a crucible of itself and this minimizes reactions with the crucible. This allows the preparation of high-purity films. The refractory oxides undergo chemical fragmentation during evaporation resulting in an undesired stoichiometry. Hence, reactive evaporation is used in which the vapors of the evaporated material are carried by the reactive gas (oxygen, in this case), and when the thermodynamic conditions are met, the vapors react with the gas in the vicinity of the substrate to form films of required composition. During the deposition of the film, the optical thickness is monitored and controlled by means of an optical monitoring system and a quartz crystal monitor. The deposited films were characterized for optical properties using a UV-VIS-NTR scanning spectrophotometer in the wavelength domain of 190 nm - 1200 nm. The experimental data so obtained were fitted to obtain optical constants, namely, refractive index («) and extinction coefficient (k), using the Sellmeier-Urbach dispersion relations given in equations (2) and (3) below [5,6]:

niE) = fÎS (2)

k(A) = aexpß 124001 - - - j (3)

where Ed (dispersion energy), E„ (Oscillation energy), a, ß, and ^are the model parameters. The films were also characterized for surface topography and microstructure using atomic force microscope (AFM). The AFM, invented by Binnig and Gerber in 1986 [7], is an important tool for surface characterization. Contact mode was used in this study as this is suitable for hard surfaces and provides high resolution. In this mode the steric repulsive force is utilized and the tip is very close to the specimen surface. The surface morphology of thin film samples conventionally describes the root-mean-square (RMS) roughness, which is a parameter on which light scattering primarily depends. However, for a systematic study, it is better to deal with a Fourier-transformed spectral surface roughness parameter called power spectral density (PSD) function that provides valuable information on height deviation on roughness profile as well as the lateral or spatial distribution of heights [8]. PSD functions describe two aspects of the surface roughness such as the spread of heights from a mean plane, and the lateral distance over which the height variation occurs [9]. The computation of two dimensional PSD function adopted in this work is given by [10]:

where S2 denotes the two-dimensional PSD, L2 is the scanned surface area, N is the number of data points per line and row, Z„„ is the profile height at position (m, ri),fx,fy are the spatial frequency in the x- and j-directions and AL=N/L is the sampling distance.

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The PSD function for the ZKVMgO as well as ZrCh-MgO-AhCb composite films was computed from AFM morphological analysis. After being characterized for optical properties and surface roughness, the optimized binary and ternary oxide composite thin films were used for the synthesis of a high-reflectance mirror for a design wavelength of 550 nm. Reactive electron beam physical vapor deposition under the optimized conditions was followed by characterization using spectrophotometer for optical properties and AFM for surface characteristics. Also, Grazing Incidence X-Ray Reflectivity (GIXR) technique was used to obtain important data like interfacial roughness as well as density measurements.

Results and Discussion

The transmittance data for the single layer thin films experimentally obtained and fitted using the dispersion relations given in equations (2) and (3) yielded the variation of refractive index versus wavelength (Figures 1 and 2). Deposition was carried out at oxygen partial pressures of 8xl0"5 mbar, lxlO"4 mbar, 4X10"4 mbar, and ôxlO"4 mbar, as well as substrate temperatures of300°C and 150°C.

Figure 1. Dispersion data for binary oxide composite thin film deposited at 300°C and various oxygen partial pressures.

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Figure 2. Dispersion data for ternary oxide composite thin film deposited at 300°C and various oxygen partial pressures.

The thin films have been characterized for surface topography and microstructure using AFM. Some of the images obtained are given in Figures 3 and 4.

Figure 3. AFM images of Zr02-MgO composite film, T = 300°C, p = 8xl0"5 mbar.

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Figure 4. AFM images of Zr02-MgO-Al203 composite film, T = 300°C, p = 8xl0"5 mbar.

The PSD function for the ZrCte-MgO as well as ZrCte-MgO-AkCb composite films has been computed from AFM morphological analysis (Figures 5 (a) and (b)).

Figure 5. PSD function for (a) binary oxide and (b) ternary oxide composite films deposited at various O2 partial pressures, T = 300°C

It is clearly seen that optimization of composite thin films to be used for a high reflection mirror application is a compromise between the surface roughness and the refractive index. As already discussed, to obtain a narrow band HR mirror, the difference between the refractive indices of the high- and low-index materials, in this case the binary and ternary oxide composites, should be minimum. Although the roughness of a film may be higher, leading to greater scattering, it may be chosen if under the same conditions of deposition, both binary and ternary oxides have very close refractive indices. Hence, the binary and ternary thin films deposited at 300°C at a rate of 5 A/s under an oxygen partial pressure of 8x10" mbar, have been found to be optimum for the purpose of developing a narrow band HR mirror for a design wavelength of 550 nm. They have refractive indices 1.76 and 1.73 respectively at this wavelength. The high-reflection laser mirror developed by depositing 51 alternate layers of binary and ternary oxide composites showed about 99.7% reflectance at the Nd:YAG second harmonic laser wavelength of 532 nm, having a narrow bandwidth of 60 nm. Transmission data has been obtained at normal incidence and at an angle of 45°. The results are shown in Figure 6 below.

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Figure 6. Reflectance data for the HR laser mirror

The HR mirror has been characterized using GIXR technique and important data like surface and interfacial roughness values as well as density measurements have been obtained. The experimental and fitted data is shown in Figure 7. The binary oxide composite is seen to have a higher density of 5.007 g/cm3 as compared to 4.203 g/cm3 for the ternary oxide composite. The surface has a high roughness of 24.29 Â and the interfacial roughness between the top layer and the one below it is 19.99 Â.

Figure 7. Data obtained from Grazing Incidence X-Ray Reflectivity technique

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Conclusion

Reactive electron beam physical vapor deposition process has been used to prepare single layer thin films of binary ZrÜ2-MgO and ternary Zr02-MgO-Al203 composites. These films have been characterized for optical properties and morphology by spectrophotometer and atomic force microscopy. The optimized composite thin films with appropriate refractive index and surface quality have been used to develop a designated high-reflection (HR) laser mirror utilizing a periodic sequential deposition of alternate high- and low-index oxide composites. This sequentially deposited design is found to be very stable and the compatibility of the two materials is demonstrated by the fact that no peeling off (instability) is observed in this 51-layer stacked system. GIXR characterization has yielded interesting data pertaining to the HR mirror, viz., roughness, interface diffusion and density values, etc. The quality HR mirror developed has potential use in laser and spectroscopic applications especially, involving high energy and high power pulse and CW laser radiations.

References

[1] N. K. Sahoo and A. P. Shapiro, "Process-Parameter-Dependent Optical and Structural Properties of Zr02MgO Mixed-Composite Films Evaporated from the Solid Solution,"^/;/. Opt., 37 (1998), 698-718.

[2] Naba K. Sahoo and Alan P. Shapiro, "MgO-Al203-Zr02 Amorphous Ternary Composite: A Dense and Stable Optical Coating," Appl. Opt., 37 (1998), 8043-8056.

[3] Milton Ohring, The Materials Science of Thin Films (San Diego, CA: Academic Press, 1992), 513.

[4] P. H. Lissberger, "Optical Applications of Dielectric Thin Films," Rep. Prog. Phys., 33 (1970), 197.

[5] N. K. Sahoo et al., "Morphological, microstructural and optical properties supremacy of binary composite films—A study based on Gd203/Si02 system," Applied Surface Science, 253 (2007), 3455-3463.

[6] F. Atay et al., "Optical characterization of Sn02:F films by spectroscopic ellipsometry," Journal of Non-Crystalline Solids, 356 (2010), 2192-2197.

[7] G. Binnig, C. F. Quate, and Ch. Gerber, "Atomic Force Microscope," Phys. Rev. Lett, 56 (1986), 930-933.

[8] R. B. Tokas et al., "A comparative morphological study of electron beam co-deposited binary optical thin films of Hf02:Si02 and Zr02:Si02," Current Applied Physics, 8 (2008), 589-602.

[9] Anthony J. Perry, "The surface topography of titanium nitride made by chemical vapor deposition," Surface and Coatings Technology, 132 (2000), 21-25.

[10] Josep Ferré-Borrull, Angela Duparre, and Etienne Quesnel, "Procedure to Characterize Microroughness of Optical Thin Films: Application to Ion-Beam-Sputtered Vacuum-Ultraviolet Coatings," Appl. Opt., 40 (2001), 2190-2199.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Microstructure and Wear Properties of Laser in-site Formation of TiBx and TiC Titanium Composite Coatings

Jing Liang ', Changsheng Liu ', Suiyuan Chen', Cuixia Ren '

Key Laboratory for Anisotropy and Texture of Materials of Ministry of Education, Northeastern University, Shenyang, Liaoning, 110004, China

Keywords: Laser in-situ formation, titanium composite coatings, Ti-6A1-4V, TiBx, TiC

Abstract

Ti-6Al-4V/B4C mixed powder which are pre-pasted or synchronized fed on Ti-6A1-4V substrates separately were scanned by a 500W pulsed YAG laser to in situ format titanium composite coatings contained TiBx and TiC ceramic reinforced phases. The influences of processing parameters and powder proportions on the microstructure and properties of the coatings were investigated. Three-layer composite graded coatings were formatted with prepasted B4C increased 10 wt. % each layer. Microstructures and phases of coatings were analyzed by OM, SEM, TEM and XRD respectively. Results show that two and more kinds of ceramic reinforcements were in situ formatted in the matrix of Ti-6A1-4V. The microhardness and the amount of TiB2 of the layer increased with the increase of B4C addition. The average micro-hardness of a laser cladding composite graded surface layer is up to 1050HV, which is nearly 3 times ofthat of the substrate (340HV), and the wear weight loss decreased over 4 times.

Introduction

Titanium and its alloys are characterized by low density, high specific mechanical strength, heat resistance, corrosion resistance and excellent biocompatibility which are attractive for a wide range of applications in aerospace, transportation and medical fields etc.. However, their applications are limited by poor oxidation resistance at high temperature, low heat conductivity coefficient, low wear resistance etc. l1,2'. So increasing the surface properties of titanium key parts becomes research focus recently, many of them focus on adding the ceramic particles such as TiN, TiC,TiB2,WC,SiC in the laser cladding coatings12'31. Laser scanning induced in-situ formation of two and more hard ceramic reinforced phases in the surface coating directly, which using high energy density laser beam melting the mixture powders of the composite elements such as titanium and its alloys, C, B and B4C powders etc. leads to rapid melting and solidification on the substrate gained relative 'ductile', high hardness and good wear and corrosion resistance coating on titanium alloys. The advantages are the laser induce in-situ reaction can form ceramic phases with different compositions, the size of the reinforcements are relative small with even distribution, strong bonding with the bonding matrix formed by laser scanning the metal powder. A laser in-situ formed TiB/Ti-6Al-4V coating shows a good wear resistance141. B J Kooi et.al. w investigated the evolution of the microstructure of TiB/Ti coatings using laser cladding Ti/TiB2 mixed powder on the Ti-6A1-4V substrate. According to three morphologies of the TiB obtained in the laser scanning layers, their growth mechanisms were indicated. Banerjee e?.a/.'5'. used direct laser fabrication to prepared Ti-6Al-4/TiB composite materials by adding Ti-6Al-4Vand 2% B element powders and investigated the structure of the in-situ prepared TiB dispersed particles. F Wang et. al. t6'7' using Ti-6AMV powder or wire together with the TiC powder direct laser fabricated the functional graded

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materials and also investigated the processing parameters and the wear resistance of the Ti-6A1-4V/TiB composite material formed by laser scanning inject pre-mixed TiB2/Ti-6Al-4Vpowder. D. Liu ef. a/.'^investigated the microstructure, tensile properties and creep rupture behaviors of the laser melting deposited in-situ TiC/TA15 titanium based composites. Recently more researches investigated on laser scanning with the addition of B4C powder to induced in-situ formation of titanium composite coatings or materials. H Tian et.al.lm and Y J Liang et.al.lu]

both using B4C/Ti mixed powder to prepared in-situ formation of TiB/TiC reinforced phases on the Ti-6A1-4V substrate. X.T.Yu et.al.[ 131 prepared TiB/TiC reinforced composite materials by laser melting deposition of B4C/TAI5 powders. In present work, T1-6AI-4V/B4C mixed powder was used as a laser cladding material prepasted or synchronized fed on a Ti-6A1-4V substrate which scanned by a 500W pulsed YAG laser to induced in situ formation of a titanium composite coating contained TiBx and TiC ceramic reinforced phases.

Experimental Procedures

B4C powder bought from Mudanjiang Qianjin Boron Carbide Ltd. with size about 50~150^m, purity no less than 90% with impurity of C, Ti-6A1-4V powder prepared by HDH (Baoji Maite Ltd. with the size ranging from 109 to 250um, purity of 99% ) for the powder prepasted specimens and that prepared by PREP (Xi'an Baode Powder Ltd. with size about 106~150um ) for the powder injected ones were used in the experiments respectively. Ti-6Al-4V/B4C with different composite ratio 10B4C+90Ti-6Al-4V(wt.%), 20B4C+80Ti-6Al-4V(wt.%) 30B4C+70Ti-6Al-4V(wt.%) were prepasted on Ti-6A1-4V substrates(17*8><12 mm3). Three-layer composite graded coatings were formatted with prepasted B4C increased 10 wt. % each layer. The thickness of prepasted powder layer was 0.4±0.02 mm. 10B4C+90Ti-6Al-4V(wt.%) was synchronized fed to prepared coatings compared with the prepasted ones.

A JHM-1GY-400 Nd:YAG pulsed laser was used. The laser processing parameters were as follow: pulse width (PW): 1-3.0ms, pulse frequency (PF):10~40Hz, the laser energy of single pulse (Es): 0.6-5.6J, and the laser scanning speed (V): 0.3-4.5mm/s. For the laser cladding with synchronized powder feed, the processing parameters also include the Ar gas flow rate of 8L/min and the powder feed rate (PFR): 4.5-2lg/min. Ar gas protection was used during the processing to avoid oxidizing. The composition and the microstructure of the cross-section of laser cladding layers were characterized by an 01ympus-GX71 optical microscopy (OM) and a Shimadzu SSX-550 scanning electron microscope (SEM). TEM specimens of the laser cladding layers were analyzed by a TECNAI G220 operating at 200kV. X-ray diffraction (XRD) analysis was carried out on a Philips X'pert PW3040/60 diffractometer with Cu Ka radiation of wavelength 1.5406 Â operated at a voltage of 40kV, a current of 40mA and a scanning rate of 107min (28=20-100°).The microhardness of the cross-section of the laser cladding layers was measured by a Wilson Wolpert 401MVD digital microhardness tester with a load of 50g and a holding time of 10 seconds, each data presented was the average of 3 measurements. Slide wear test was performed using a MG-2000 type wear test machine (pin-on-disc) with a load 50N, rotate rate of 250r/min, the distance of 122.5m, 245m and 367.5m. The specimens were weighted on an electrical balance with the accuracy of O.OOOlg.

Results and Discussion

Though a series of experiments, initial laser processing parameters were set: PW=3.0ms, PF=15 Hz. The defocused length (DL) was 15mm for the powder prepasted specimens. The laser power and laser scanning speed were optimized with the above processing parameters keeping in

28

constant. When the laser line energy in the range of about ll~12J/mm for laser cladding prepasted Ti-6Al-4V/B4C and 18.6J/mm for powder inject 10B4C+90Ti-6Al-4V( wt.%), single path specimens were produced with the good metallurgical bonding with the substrates, and the distribution of the in-situ formatted ceramic phases was even in the mass. Multi-path powder prepasted coatings were formed with one of the optimized processing parameters (Es: 2.4J, V: 3.0mm/s, overlapping rate of 50%). The optimized processing parameters of the powder feed laser cladding specimens in the case are as follows: DL: 20mm, Es: 5.6J and V: 0.3 mm/s together with Ar gas flow rate of 8L/min and the powder feed rate of 9.97g/min accordingly.

The microstructure of the cross-section of the titanium composite coatings with powder prepasted was shown in Fig 1. Two and more kinds of ceramic reinforcements with different morphologies were in-situ formatted in the matrix of Ti-6A1-4V. The overall distribution of the ceramic phases was relatively even and no crack and prose was found in the coatings. The tiny dendrites with all directions microstructure is throughout the coating while some of them near the coating surface had orientation (see Fig 1 and Fig 2c), that the growth of the dendrites was restrained by each other along the vertical direction to the temperature gradient due to the high cooling rate of the laser scanning process from the substrate to the top of coatings. Along the converse direction to the temperature gradient, the grain rapidly grows to form dendrites due to the transient exist of the liquid metal, the growth direction of the dendrites of the non-interface areas inside the coating depend on the orientation of the matrix. The bonding matrix for the hard ceramic reinforced phase here was produced by laser scanning of the Ti-6AMV powder in the mixture of powder addition. Relative "ductile", high microhardness and good wear resistance titanium composite coatings were thus gained. Some needle-like or fiber-like tiny ceramic phases were also found in the coatings (see Fig 2a) according to the literature maybe TiB|5121 which confirmed by TEM results. Because the TiC possesses a NaCl-type structure that make a TiC tend to grow into a spherical form to reduce the interfacial energy ' , so small disperse particles found in the coating maybe TiC (see Fig 2a), while in the relatively high temperature molten pool of laser scanning, TiC also had time to grow up to relatively large dendrites (see Fig 2b). For laser cladding prepasted 30B4C+70Ti-6Al-4V (wt.%) powder mixture, some of TiB2(hexagonal prism) were found(see Fig 2d) and TiC dendrites size decreased because of B element is rich in some area with the increase of B4C amount in the layer. During the solidification, the solubility of B in the beta titanium phase is large difference with that in the alloy liquid. This offers additional constitutional supercooling at the solidification front that increased the nuclear rate of the TiC. Moreover B is active in the high temperature which is easy attached to the surface of the TiC. This prevents TiC to grow up and makes TiC second dendrites form arms with small diameter which melt to be short bars or disperse particles (see Fig 2c).

Fig. 1 OM morphology of the single path powder prepasted laser cladding coatings, (a) 10B4C+90Ti-6Al-4V, (b) 20B4C+80Ti-6Al-4V (c) 30B4C+70Ti-6Al-4V ( all in wt%).

29

Fig.2 (a~d) SEM morphology and (f) XRD analysis of the single path powder prepasted laser cladding coatings( all in wt.%),

(a)10B4C+90Ti-6Al-4V, (b) 20B4C+80Ti-6Al-4V (c, d) 30B4C+70Ti-6Al-4V.

The X-ray diffraction results of single path laser cladding specimens with different prepasted powder composition were shown in Fig 3.Together with the former microstructure analysis, it can be seen that in the laser cladding 10B4C+90Ti-6Al-4V and 20B4C+80Ti-6Al-4V (in wt.%) specimens, the coatings contained a large amount of TiB and TiC (Fig 2a,2b). The chemical reactions in the laser scanning produced molten pools were as follows.

5Ti+B4C=4TiB+TiC (1)

3Ti+B4C=2TiB2+TiC (2)

Researches show that AG for reaction (1) and (2) are negative, and the reaction AH are large.The two reactions are exothermic reactions which spontaneously sustained [14'. So throughout the temperature range of 500K to 2500K, TiB2 and TiB both are probable reinforcement particles. TiC reinforcement can be preferentially synthesized by the two reactions. However, TiB2 is not stable when an excess of titanium exists in the Ti-TiB2 system and can be transformed to TiB according to the reaction (3) as follows.

TiB2+Ti=2TiB (3)

AG of reaction (3) is also negative and the reaction occurs thermodynamically as well. The titanium existed in the molten pool produced by laser scanning prepasted 10B4C+90Ti-6Al-4V and 20B4C+80Ti-6Al-4V (in wt.%) powder mixture was excess, and the cooling rate in laser cladding prepasted powder is relatively slower than that in the laser cladding with inject powder. So Ti element had time to diffuse which made high concentration near die B4C particles, so large amount of TiB was formed in this situation. While for the laser cladding prepasted 30B4C+70Ti-6A1-4V (in wt.%) specimens, the B element is much more than the former one with the increase of B4C addition, when the Ti element is consumed that reaction (3) can't happened, so some TiB2 formed in the high temperature is left without being transformed into TiB in this situation.

For the laser cladding with synchronized powder feed specimens, the cross-section microstructure was quite different from that of the powder prepasted coating. More TiB2

(hexagonal prism or rod-like phase) reinforced phase appeared in the centre of coating (see Fig

30

4b) while dendrites and disperse particles near the interlace. The melting points of TiB2 and TiC is about 3200"C[I0], so during the solidification, TiB2 will first nucleated and grow to form prism and at the same time TiC phase will nucleate and grow as well. Because the melting point of TiB is about 2200'C, which is much lower than that of T1B2, so during the solidification of laser cladding, T1B2 predominately nucleated and had time to grow coarser. TiB nucleated later than TiB2 and then form the tiny fiber-like morphology.

Fig. 4 Cross-sectional morphology of titanium composite layers fabricated by laser cladding 10B4C+90Ti-6Al-4V,with synchronized powder feed under the process of Es=5.6, V=0.3mm/ss

(a) OM micrograph, (b, c) SEM micrograph (deep etched)

During the laser cladding with synchronized powder feed, the react time of the laser and the powder was quite short; the time for the molten pool formed to its solidification was also short. When the Ti liquid phase contact with the B4C solid particle, it consumed rapidly and the continuous melting of the B4C made the area with high concentration of B element but low. When the reaction continued the concentration of Ti in the area is not enough, so the reaction (3) seldom occurs. On the other hand, the energy input for the laser cladding with synchronized powder feed specimen was also much higher than that for the prepasled powder one, and the laser reacted directly with the powder flow without consuming energy on the organic bond material as prepasted powder laser cladding does, together with the gas flow for the synchronized powder feed all increased the cooling rate of the laser cladding process at this condition.

Fig 5. XRD patterns of laser in-situ fabricated single-path titanium composite layers with synchronized powder feed.

The thickness of the three layer composite graded coating was about 150 pm (see Fig 6a) with laser cladding powder mixture of 10B4C+90Ti-6Al-4V on the substrate followed 20B4C+80Ti-6A1-4V and 30B4C+70Ti-6Al-4V(all in wt.%) on the top layer. No cracks were found in the gradient coatings even the B4C in powder mixture of top layer also reaches 30 wt. %. From the

31

EDX line scan, the amount B and C graded increased while the amount of Ti decreased with the increase of B4C from the substrate to the top layer. The microstructure of the three-layer composition graded coating (Fig 6) were quite different from the single composite one (Fig 2c,d). A great amount TiB2 was found throughout the coating with some orientation near the surface (Fig 6b, 6c and Fig 7 arrow point). More TiC disperse particles were found around the tiny dendrites. During the laser cladding of the three-layer composition graded coatings, with the increase of the B4C addition in the powder mixture, the amount of B element increase that prevented the TiB2 change into TiB according the Eq.(3) and made the size of TiC finer.

Fig. 6 SEM morphologies of the top layer of the laser cladding B4C /Ti-6A1-4V gradient coating.

Fig 7 XRD analysis of the top layer of the laser cladding B4C /Ti-6A1-4V gradient coating.

The distribution of the microhardness of the cross-section of the powder prepasted laser cladding layer from the surface to the substrate was shown in Fig 8. It can be seen that the microhardness was in the tendency of graded decrease from the coating to the substrate. Because the microhardness of the evenly distribution laser induced in-situ formation of TiC (up to HV3200) and TiB (up to HV3600) ceramic phases was very high, which greatly increased the microhardness of the titanium composite coating. The maximum microhardness of the one-layer coating up to 800HV which is nearly twice that of substrate (340HV) and for the three-layer gradient coating it reaches 1050HV which is 3 times that of the substrate. The thickness of the

32

three-layer gradient coating indicated in Fig.8 was about 0.15 mm agreed with that shown in the OM micrograph (see Fig6a). The improvement of the microhardness of the HAZ was due to the formation of the needlelike martensite in this area.

Fig 8 Microhardness results of the laser cladding coatings.

The results of the pin-on-disc experiments have shown in Fig 9. The multi-path powder prepasted specimens for the wear test were prepared under the processing conditions: ES=2.4J, V=3.0mm/s and overlapping rate 50%. It is shown that the sliding wear property of the titanium composite coating were greatly improved compared with that of the substrate because of the in-situ formation of the TiB and TiC hard and wear-resistant reinforced phases. The wear loss is the specimen weight loss during the wear tests. When the sliding distance is 122.5m, the average wear loss of the substrate was 0.0082g, while that of the one-layer 20B4C+80Ti-6Al-4V (wt. %) coating for comparison was only 0.0012g. With the increase of the sliding distance to 245m, the average wear loss of the one-layer 20B4C+80Ti-6Al-4V (wt.%) coating increased to 0.0023g which is higher than that of the three-layer single composition 30B4C+70Ti-6Al-4V (wt.%) specimen (0.0017g). The average wear loss of the three-layer composition graded specimen under the longer sliding distance of 367.5m is only 0.0019g which is much lower than that of the three-layer single composition 30B4C+70Ti-6Al-4V (wt.%) specimen(0.0050g). The wear loss of the gradient titanium composite coating decreased relatively over 4 times than that of the substrate. The friction coefficient of all the multi-path coatings is about 0.2-0.3 which is also lower than that of the Ti-6A1-4V substrate (0.4).

Fig. 9 The comparison of the wear weight loss of the laser cladding coatings.

33

Summary

Titanium composite coatings contained TiBx and TiC ceramic reinforced phases in the matrix of Ti-6A1-4V have been successfully in situ formed under optimized processing parameters by laser scanning Ti-6Al-4V/B4C mixed powder which was either prepasted or synchronized fed on a Ti-6A1-4V substrate. TiB and TiC ceramics were formed evenly with the morphology of needle, tiny dendrites and disperse particles in the prepasted single path specimens and a small amount of TiB2 in the 30B4C+70Ti-6Al-4V (wt.%) powder prepasted specimens. The microhardness and the amount of TiB2 of the layer increased with the increase of B4C addition. For the synchronized powder feed laser cladding layer, the ceramic reinforcements were TiB2 (hexagonal prism and rodlike), TiB (needlelike), a small amount of TiC (disperse particles) and non fully reacted B4C. The average microhardness of a laser cladding composite graded surface layer is up to 1050HV, which is over 3 times ofthat of the substrate (340HV), and the wear weight loss decreased over 4 times.

Acknowledgements

The author gratefully acknowledges Scientific Research Foundation (SRF) from State Education Ministry and Northeastern University, and Support by Program for Changjiang Scholars and Innovative Research Team in University (IRT0713), respectively.

References

1 B. J. Kooi, Y. T. Pei, and J Th M. Hosson, "The Evolution of microstructure in a laser clad TiB-Ti composite coating" . Acta Materialia, 2003, 51: 831~845.

2 Y. T. Pei, V. Ocelik, and J Th M. Hosson, "SiCp/Ti6A14v functionally graded materials produced by laser melt injection" . Acta Materialia, 2000,50: 2035-2051.

3 J. A. Vreeling et al. ,"Ti-6Al-4V strengthened by laser melt injection of WCP". Acta Materrialia, 2002, 50: 4913-4924.

4 V. Ocelik, D. Matthews, and J Th M. Hosson, " Sliding wear resistance of metal matrix composite layer prepared by high power laser". Surface and Coating Technology, 2005, 197(2-3): 303-315.

5 R. Banerjee et al., "Direct laser deposition of in situ Ti-6Al-4V-TiB composites". Material Science Engineering, 2003, A358: 343-349.

6 F.Wang, J.Mei, and X. Wu, "Direct laser fabrication of Ti6A14V/TiB". Materials Processing Technology. 2008, 195: 321-326.

7 F.Wang et al., "Laser fabrication of Ti6A14V/TiC composites using simultaneous powder and wire feed". Material Science Engineering 2007, A445-446: 461-466.

8 D Liu, S.Q. Zhang, A. Li, H.M Wang. Alloy Compound 2009, Vol.485:156-162. 9 D Liu, S.Q. Zhang, A. Li, H.M Wang. Materialsand Design, 2010, Vol.31:3127-3133. 10 H. Tian et al., "Microstructure of laser cladding B4C and B4C+Ti coatings on TC4 alloys". Rare Metal

Materials and Engineering, 2007, 36(3):420-423. 11 Y. J. Liang et al., " Laser cladding deposition on in-situ formation of titanium composite materials".

Nonferrous Metals, 2008, 60:25-28. 12 X.T. Yu and H. M. Wang, "Microstructure of laser melting deposited (TiB+TiC)/TA15 titanium

matrix composite bar". Aerospace material and technology, 2007,6:116-119(in Chinese). 13 X.T. Yu and H. M. Wang, "Microstructure and mechanical properties of laser melting deposited

(TiB+TiC)/TA15 in situ titanium matrix composites". Acta Materiae Compositae Sinica, 2008, 25:113-118 (in Chinese).

14 W. J. Lu et al., " Solidification paths and reinforcement morphologies in melt-processed (TiB+TiC)/Ti in situ composites". Metall. Mater. Trans A, 2002, 33: 3055-3066.

34

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Creep behavior of plasma sprayed Y-PSZ coated 6063-T6 aluminum alloy

Eray Erzi, Cem Kahruman, Suat Yilmaz

Metallurgical and Materials Engineering Department, Istanbul University 34320, Avcilar, Istanbul, Turkey

Keywords: Plasma spraying, Y-PSZ coating, 6063 aluminum alloy, creep.

Abstract

Aluminum alloys have been used in industries for years because of light weight. But these materials have some disadvantages like low mechanical properties. For increasing of the mechanical properties of aluminum alloys can be coated with hard ceramics. In this study, T6 heat treated 6063 aluminum alloy were coated with %8 Y-PSZ powders by plasma spraying method. Subsequently, creep properties of the uncoated and coated samples were investigated. Temperature dependent creep tests were carried out at 100, 125, 150°C temperatures and at 170 MPa stress conditions. Stress dependent creep tests were carried out at 120, 145, 170 MPa stresses and at 100°C temperature conditions. The creep results of the uncoated and coated samples were discussed and compared.

Introduction

Zirconia (ZrC>2) is a ceramic material with important characteristics. Low thermal conductivity, good corrosion resistance, high fracture toughness and its transformation toughness effect are some of these characteristics. Due to these properties zirconia is used as a thermal barrier coating and wear resistant material [1-3]. In this study 6063-T6 aluminum alloy is coated with yittria stabilized zirconia by plasma spray method and the effect of this coating investigated via creep tests. Time dependant deformation under a constant temperature and load that occur in material is called creep. During the creep test the deformation rate is measured as a function of time and the creep curve is obtained as shown in Figure 1.

Figure 1. a)T:constant (o i >a 2 >o 3 ) , (é i>£2>£3) b) o xonstant(Ti>T2>T3),(É 1 > é 2 >£ 3)

35

As the figure shows the primary creep (primary), secondary (steady creep) and tertiary (third) creep curve are divided into three regions. Dislocations move rapidly through the first stage so that sample deformations rapidly. Later, dislocations flow rate slows down because of irregularities. In the second stage stable creep zone the deformation rate stabilize. In this zone the slope of the creep curve gives the minimum creep rate (é = ds/dt). Due to the high rates of deformation in the first and third regions second region is taken in to account in the engineering design. The total amount of creep deformation is (s), time (t), temperature (T) and stress dependent (é = f(e, t, a, T) [4, 5].

Experimental

The composition of the samples which were used in the experiment is 6063 aluminum alloy consist of %0.40 Mg, %0.40 Si and %0.10 Cu approximately and it can be age hardened. Used samples were produced from 30mm rod which was produced with injection press. These rods then exposed to T6 aging process. Rods first heated up to 480 C then heat treated at mis temperature for 20 minutes and then water quenched to room temperature. Then rods heated to 180°C and heat treated for 7.5 hours for aging. Heat treatments were done with Lenton UAF 15/10 model laboratory furnace. Aim of this heat treatment process is to gain of better machinability.

Y-PSZ ceramic powders (Y2O3) which were used as coating material were Powder Alloy Corp. (USA) product. 6063-T6 heat treated samples first surface roughened for better grip in the coating stage. Later a lining material was coated on the 6063-T6 aluminum base. The purpose of this lining is to bind the metal base and the ceramic coating. The lining material was coated by Sulzer-Metco 3B model plasma spray gun and the lining material was Metco Amdry 956 product and has an average size of ~45um. After lining; Powder Alloy Corp. (USA) product ~50um grain sized yttria (Y203) stabilized zirconia (Zr02) powders were coated with the same plasma spray gun (Figure 2).

Figure 2. Standard test samples for tensile creep experiment (uncoated and Y-PSZ coated 6063-T6)

36

Materials and Method

Creep test device is a tensile type MA YES brand TC-30 model (UK). Working principle of creep test device is given in Figure 3.

Figure 3. Working principle of the creep test device.

Temperature measurements made with a pair of K-type thermal element and the temperature values are taken on the sample. To investigate temperature dependent creep properties ai-Ti, C1-T2, CT1-T3 (Ti>T2>T3 with fixed stress); and to investigate stress dependent creep properties CT1-T1, 02-Ti ve C73-T1 (ai>a2>cJ3 with fixed temperature) experiment sets are created. 10 different creep tests were done for both uncoated and Y-PSZ coated 6063 aluminum samples Applied stresses respectively u\ = 170MPa, a2 = 145MPa and a3 = 120MPa were selected. Temperatures selected for experiments are T,= 150°C, T2 = 125°C and T3 = 100°C. Experiment sets are given in the Table 1 according to the generated temperature and stress dependent creep test parameters.

Table 1. Selected and used stress and temperature values in the creep tests.

Applied Temperature, Tapp.

Applied stress, ovp. T l = i5 00c T2=125°C T3=100°C

a, = 170 MPa X X X o2= 145 MPa X g3 =120 MPa X -- --

Determination of the creep properties of Y-PSZ coated and uncoated 6063 aluminum alloys , calculations of creep rate (£) and effect of Y-PSZ zirconia coating to the creep behavior to 6063 aluminum alloy under constant temperature and constant stresses.

37

Results

Constant stress and different temperature (CTI - Tl, CTI - T2, ai - T3) conditions comparative test results for 6063-T6 aluminum alloy are given in Figure 4a-b (Fig. 4.a is uncoated and 4.b is coated 6063-T6 aluminum alloys creep test results).

Figure 4 a and b. Creep behavior of samples at 170 MPa constant load and at different temperatures (CTI - Tl, oi - T2, c?i - T3) (A: uncoated, B: coated)

Constant temperature and different stress conditions (ai - Tl, en - T2, ai - T) comparative test results are given in Figure 5a-b (Figure 5a is uncoated and 5b is coated 6063 aluminum alloys creep test results).

Figure 5 a and b. Creep behavior of samples at 150 °C constant temperature and at different stresses (ai - Tl, a2 - Tl, 03 - Tl) (A: uncoated, B: coated)

Creep rates are calculated from the creep curves in Figures 4-5. The creep rates for;

38

Experiments uncoated (mm/h) coated (mm/h)

170MPa-100°C -» 0.00106 0.00017 170MPa-125°C -> 0.025 0.032 170MPa-150°C -> 2.0 0.9 i45MPa-150°C -» 0.014 0.00152 120MPa-150°C -» 0.000148 0.000321

Conclusion

With increasing temperatures and stresses both uncoated and coated samples creep rates increases. However uncoated samples creep rate is higher and rupture time is shorter than Y-PSZ coated samples. Y-PSZ coated 6063 aluminum alloy has higher creep resistance.

The creep values are important for materials selection and CAD/CAE modeling.

Acknowledgement

This work was supported by Research Fund of Istanbul University (Project Nr.: 515/05052006).

References

1 P. Ramaswamy, "Evaluation of Cao-Ceo2-Partially Stabilized Zirconia Thermal Barrier Coatings, "Ceramics International, 25 (1999), 317-324.

2 J.H. Shin, "Effect of Annealing and Fe203 Addition On The High Temperature Tribological Behavior of The Plasma Sprayed Yttria-Stabilized Zirconia Coating," Surface andCoatings Technology, 133-134 (2000), 403-410.

3 X.Q. Cao, "Ceramics Material for Thermal Barrier Coatings," Journal of the European Ceramic Society, 24 (2003),1-10.

4 S.P. Deshmukh, et al., "Creep Behavior and Threshold Stress of an Extruded Al-6Mg-2Sc-lZr Alloy," Materials Science and Engineering A, 381 (2004), 381-385

5 J. CADEK, et al "Threshold creep behaviour of aluminium dispersion strengthened by fine alumina particles," Materials Science and Engineering: A, 252 (1) (1998), 1-5

39

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Contribution of Ti Addition to the Electronic Structure and

Adhesion at the Fe2Al5/Fe Interface in 55%A1-Zn Coating

Guangxin Wu1'" * Yuling Ren2) Jieyu Zhang1*' * Kuochih Chou1' 1. School of Materials Science and Technology, Shanghai University, Shanghai 200072, P.R. China

2. Cold Rolling Plant, Baosteel Branch, Baoshan Iron & Steel Co. Ltd., Shanghai 200941, P.R. China

Abstract

We report a density functional theory investigation of the atomic structure, bonding, and ideal work of adhesion of the Ti addition to Fe2Al5/Fe interface, in order to explore the potential of Al-Zn-Si-Ti as a protective coating for steel. The results show that the clean interface has an ideal work of adhesion of 13.92 eV and 12.35 eV for Fe-terminated and Al-terminated interface while the presence of Ti increase adhesion to 15.01 eV and 13.64 eV via rearrangement of electron density upon formation of the interface. Then, the fracture toughness of Ti dopant Fe2Als/Fe interface are discussed with Griffith fracture theory and results show that Ti addition decreases embrittling effect at Fe2Als/Fe interface. Our results provide theoretical evidence for the excellent adhesion behaviors of Ti addition Fe2Als/Fe interface. Keywords: First-principles; Al-Zn-Si coating; Electron structure; Griffith fracture work

Introduction

The 55%Al-Zn-1.6%Si coated steel product (GALVALUM) is one of the popular alloys in the building industries, and has distinct advantages by providing not only an better anticorrosive performance than that of zinc coating but also a high temperature oxidation resistant the same as aluminum coating | l j . Therefore, these alloys have attracted extensive interest worldwide.

The presence of alloying elements in the bath has been shown to affect strongly the morphology, kinetics and growth of the zinc-alumum coatings during the hot-dip process. The addition of Ti to this known composition of 55% Al-Zn-Si coating has been proposed in several patents and literatures for a number of years [2-4]. It has been show that Ti additions to the liquid Al-Zn-Si bath would enhance the nucleation rate when solidification takes place on the surface of the strip as it leaves the pot. This effect would cause then an additional refining effect on the spangle size of the strip and improve its surface properties[5-8J.

In modern continuous hot-dip galvanizing line, the steel strip travels at speeds between 100 and 200 m/min through a ceramic pot filled with liquid Al-Zn-Si alloy held at approximately 600 °C 11|. Then, the rapid exothermic reaction between the Zn-Al-Si coating and the steel substrate will take place during the short time period by

* Corresponding author. Tel.: +86 021 56337920; fax.: +86 021 56337920

E-mail address: gxwu@shu.edu.cn (Guangxin Wu); ziv6l62@staff.shu.edu.cn (Jieyu Zhang)

41

promoting the formation of a thin and continuous solid reaction layer at the coating/steel interface [9]. These reactions are so rapid as to make the formation mechanisms of the Fe-Al inhibition layer difficult to study. Two orientation relationships between the Fe2Al5 and the steel substrate were reported by Guttmann [10] to be (311)FA//(110)Fe or (221)FA//(110)Fe, where FA represents Fe2Al5. Meanwhile, an orientation relationship of (001)FA//(011)Fe and [I30]FA//[IOO]Fe

between Fe2Al5 and the steel substrate was also reported [11]. Recently, the results of TEM electron diffraction showed that only Fe2Als, with some Zn dissolved in it, was formed and its orientation relationship with the Fe substrate was: [11 0]FA//[ 111 ]Fe, (001)FA//(011)Fe, (lI0)FA//(2II)Fe- Furthermore, the interfaces were mostly (001)FA/(0Il)Fe which is a low energy interface with good match of lattice sites [12J.

In previous, we have investigated the effect of Ti addition on the thickness of intermetallic layer[13|. However, until recently, very little has been known about the effect of Ti addition on the adhesion of coating/substrate. Therefore, in this paper, the effect of Ti on the adhesion of coating/substrate with calculated binding energy, work of adhesion and charge transfer based on first-principles method will be predicted.

Methodology

The calculations performed in this study were done using the Vienna ab initio simulation package (VASP) [14]. This program evaluates the total energy of periodically repeating geometries based on density-functional theory and the pseudopotential approximation. In this case, the electron-ion interaction can be described by the projector augmented wave (PAW) method of Blöchl in the implementation of Kresse and Joubert [15]. The generalized gradient approximation (GGA) through the Perdew-Wang 91 (PW91) [16] functional was used for the exchange-correlation potential.

A seven-layer Fe2Als (001) surface was built, the termination can be either Fe or Al, and in order to avoid creating the dipoles in the periodic cell artificially. Similarly, we also cut Fe (1 10) surface from the bulk Fe crystal with body centered cubic structure (BCC), the thickness of Fe slab was selected so that the interior of the slab will be bulk like, we found the slab with six atomic layers was sufficient. In surface energy calculation, the slab is allowed to fully relax, and is converged with respect to slab thickness. For modeling Fe2Als/Fe interfaces, a superlattice geometry was used in which a 13 and 14 layers is placed for Al-terminated and Fe-terminated structures. In order to match the lattice geometry, the (001) surface of Fe2AU was rotated to a set of new orientations, and then the Fe slabs was placed. The orientation relationship we studied in this paper is Fe2Als(001)/Fe(011). It has been proved that it is the lowest index surfaces[12], since it is the most stable and therefore the most likely to form. The schematically structure is shown in Fig. 1. In the interface calculations, the bottom three layers are kept fixed at their bulk positions to model a semi-infinite crystalline substrate. The corresponding kinetic energy cutoffs were 400 eV for all PAW calculations. Brillouin-zone integrations employed a 5x5x1 Monkhorst-Pack grid of k-points for surface and interface models.

42

Fig. 1 Schematic representation of Ti substitution in Fe-terminated (a) and Al-terminated (b) Fe2AIs(001)/Fe(011)

interface. The larger and smaller spheres indicate Fe and Al atoms, respectively. The possible occupied positions

have been titled in this figure.

Results and discussion

1. Validation for the calculations The accuracy of the computational method used in this study has been tested

initially to describe the properties of bulk Fe2Al5. In Table 1, we provide the results for the equilibrium lattice constant and enthalpy of formation of Fe2Al5 together with corresponding experimental data and other previous calculations. Good agreement with experiment gives evidence of accuracy of the GGA and the pseudopotentials used. These characteristics make us confident to pursue the next step of our investigations, namely effect of Ti addition on adhesion of Fe2Al5(001)/Fe(011) interface.

Table 1 Calculated equilibrium lattice constant (a0, *o, co) and enthalpy of formation (AH) for bulk Fe2Al5

Fe2Al5

Reference

This work

Experiment

oo(nm)

0.749

0.766 "

Ao(nm)

0.620

0.639 *

c0(nm)

0.483

0.419"

A//(KJ/mol)

-176.6

-194.0"

a Experimental value from Refs. |17] b Experimental value from Refs. j 18]

2. Site preference In order to study the site preference of Ti at the Fe2Als(001)/Fe(011) interface, the

total energies of different interfacial models corresponding to Fig. 1 are calculated and are shown in Table 2. It can be concluded that the energies of Ti substitution at Al sites in the Fe2Als(001)/Fe(011) interface are lower compared to the energies at Fe sites. This means the most favorable sites for Ti substitution in both Fe-terminated and Al-terminated Fe2Al5(001)/Fe(011) interface are Al positions. Especially, it could be seen that energies of both position 2 of Fe-terminated and position 1 of Al-terminated interface are the most negative. Thus, in the calculation of bonding energies and electronic structures below, we mainly consider the Al-substituted configurations.

Table 2 Total energy of Fe-terminated and Al-terminated Fe2Al5(001)/Fe(011) interface

Non Ti

43

Fe-terminated

Al-terminated

1

2

3

4

1

2

-541.93

-505.26

-544.21

-548.23

-548.21

-544.25

-508.60

-508.44

3. Work of adhesion The total binding energy corresponding to the addition or removal of a Ti atom at

a certain position relative to the Fe2Als/Fe interface is given by [19]

£,ot= £(clean Fe2Al5/Fe)+£(Ti)- £(fully relaxed Fe2Al5-Ti/Fe) ( 1 )

Here "clean Fe2Als/Fe" refers to the supercell containing only the Fe2Al5, Fe layers, and no Ti; "Ti" corresponds to a single Ti atom at the center of a large supercell, and "fully relaxed Fe2Als-Ti/Fe" refers to the Fe2Als-Ti/Fe supercell. Positive binding energy means that the Fe2Al5-Ti/Fe relaxed supercell is stable with respect to dissociation into Fe2Al5/Fe and a free Ti atom.

The total binding energy can be separated into two contributions [20]. The first, which we will call the chemical binding energy £c, is a positive contribution due to chemical bonding of the Ti atom to the host Fe2Als/Fe. The second, which we will call the lattice distortion energy £a, is a negative contribution due to the structural distortion of the host lattice by the interstitial. These two contributions can be extracted from the total binding energy in a straightforward fashion.

The chemical contribution is given by

Ec = ̂ (relaxed Fe2Al5/Fe) + £(Ti)- £(fully relaxed Fe2Al5-Ti/Fe) (2)

Here "relaxed Fe2Al5/Fe" refers to a Fe2Als/Fe supercell disturbed in such a way that it has the same structure as if a Ti atom was present.

The lattice distortion energy is simply the difference

Ed = £(clean Fe2Al5/Fe)- £(relaxed Fe2Al5/Fe) (3)

We see from Table 1 that the interface position for Ti is energetically more favorable than the bulk position from both the chemical binding and lattice distortion contributions. However, the main contribution to the difference is from the lattice distortion. From this we conclude that the interface position is more favorable due to insufficient volume in the Fe2Al5 to accommodate an interstitial Ti atom.

Table 3 Binding Energy of Fe-terminated and Al-terminated Fe2 A15(001 )/Fe(011 ) interface

Energy (eV)

£tot

Fe-terminated

4.42

Al-

4.98

Ei

Fe-

8.81

Al-

8.32

Et

Fe-

-4.40

Al-

-3.34

The ideal work of adhesion W^ is a rough approximation to mechanical properties of an interface. It is defined as the interface bond energy needed (per unit area) to reversibly separate an interface into two free surfaces [19]

2 rad=£(Fe2Al5-Ti slab)+£(Fe slab)-£(Fe2Al5-Ti/Fe) (4)

Here £(Fe2Al5-Ti/Fe) is the total energy of the interface system with and without Ti atom, £(Fe2Al5-Ti slab) and £(Fe slab) are the total energies of the separated slabs

44

generated by either removing the Fe2Al5 or the Fe block from the Fe2Al5-Ti/Fe interface, respectively. The factor 2 accounts for the two identical interface in the supercell structure. The W^ values for Fe- and Al-terminated Fe2Al5/Fe interface structures are summarized in Table 4. It can be seen that the Fe-terminated interfaces exhibit larger W^ values than Al- interface. The value of Wài for the Ti addition Fe-terminated interface is largest among the four structures. This means, for the two types of interfaces, the Fe-terminated structure yields the strongest interfacial adhesion. Therefore, Ti will segregate into the interface between the matrix and the precipitated phase in the metallurgical reaction could trap the Ti atom.

Table 4 The work of adhesion (W^) for the Fe2Al5/Fe interface structures

Clean

Ti-doped

Fe-terminated (eV)

13.92

15.01

Al-terminated (eV)

12.35

13.64

4. Segregation energy and Griffith fracture work Before investigating the effect of Ti on the interface, the segregation energy of a

Ti atom was calculated to assess whether a Ti atom will segregate energetically from the Fe2Al5 bulk to the Fe2Als/Fe interface. The segregation energy is the energy needed for a Ti atom to diffuse from a bulk site to the interface site and can be calculated using [21J

\pTi plot _ plot . . . L"~■Interface ^Ti-Interface ^Ti-bulk \D>

where E'^_Merfilce is the total energy of the system with a Ti atom in a interface

site, and E'^'_bulk is the total energy of the system with a Ti atom in the "bulk" of the

Fe2Al5 layers away from the interface. In this work both the Fe-terminated and Al-terminated interface were considered to calculate segregation energies, and then the most energetically favorable site was identified. The segregation energies of Ti from the bulk to the interface are listed in Table 5. It could be seen that for both interface terminations the interface position has positive segregation energy. This suggests that in experimental systems Ti atoms should prefer to reside at the interface rather than in the crystalline bulk after annealing.

Based on the above, the Griffith work can be used to describe fracture effects of Ti addition in Al-Zn-Si hot-dip coating. From Griffith fracture theory, the work of fracture is the energy separating an interface against the atomic cohesion. The influence of Ti element segregation on the Griffith work can be estimated by the value

of AE"lerßce -AE™xlure. The segregation energies at the interface have been calculated

previously. And then calculation of segregation energies at the fracture surfaces is the same as that at the interface. The results are shown in Table 5. According to Rice-Wang model, we can conclude that Ti could enhance cohesion of Fe2Al5/Fe interface, because the calculated difference of segregation energy between interface and fracture are negative for solute Ti atom. Especially, from our calculation, the Fe-terminated and Al-terminated interface (-0.92 eV and -0.37 eV, respectively) have

45

the similar result which does not affect the final conclusion. Table 5 The segeration energy and Griffith work for Fe-terminated and Al-terminated Fe2Al5(001)/Fe(011)

interface

Segregation energy (eV)

Girffith work (eV)

Fe-terminated (eV)

A£,"

-0.93

AET;

-0.01

-0.92

Al- terminated (eV)

AEj1

-0.86

AET;

-0.49

-0.37

5. Charge distribution In this section, we discuss the results on the charge distribution for Ti addition

Fe2Als(001)/Fe(0Il) interface. The charge density difference was calculated (Fig. 2) by subtracting the valence charge density of the Fe2Als(001)/Fe(011) interface with Ti from the valence charge density of the interface in which Ti is replaced by Al but the atomic coordinates of atoms remained unchanged. It is seen from Fig. 2 that Ti has a greater valence charge density difference (arising from its 3d shell), but the accumulated charge is spatially localized around the Ti core and does not take part in interatomic bonding. And it can be determined that Ti-Al bonding has a predominantly metallic character and is very similar to the Fe-Al bonding. At the same time there is an accumulation of bonding electrons around Ti, as seen in Fig. 2. The charge accumulation is somewhat spatially diffuse, but is most pronounced along the Ti-Al bonds. This accumulation indicates that Ti-Al bonds are enhancer than the corresponding Fe-Al bonds and explains the decreasing embrittling effect of Ti in Fe2Al5(001)/Fe(011) interface. Thus for cleavage along the z direction, a larger force is required to be imposed to the Fe2Als/Fe interface, which then leads to significant increase of the Griffith work.

Fig. 2 The charge density difference for (1 1 0) planes of Fe-terminated and (4 0 0) plane

of Al-terminated interface. The blue, green and red spheres indicate Fe, Al and Ti atoms,

respectively.

Conclusion

In conclusion, first-principles method was used to investigate the effect of Ti addition on electronic structure and adhesion of Fe2Ai5(001)/Fe(011) interface in

46

hot-dip 55%A1-Zn coating. The optimized geometric configurations and total energy for the Ti substitution in Fe-terminated and Al-terminated Fe2Als/Fe interface were calculated. And results indicated that the most favorable sites for Ti substitution in both interfaces are Al positions. The result of work of adhesion show than it is energetically favorable for Ti to reside at the Fe2Als/Fe interface rather than remain in the bulk of the Fe2Als. Based on Griffith's fracture theory and our calculations for surfaces and interfaces of the related materials, we found there is Ti induced ductility effects at the Fe2Als/Fe interface. Then charge accumulation of Ti addition causes enhancer Ti-Al bonding at the interface than at the fracture surface. This strong correlation between the creep-strengthening effect of alloying element and its position within the period table needs further experimental examination.

Acknowledgments

The authors would like to thank the financial supports from National Natural Science Foundation of China (51074103 and 51104098) and the Program for Changjiang Scholars and Innovative Research Team in University (IRT0739). Computational resources were supported by Asso. Prof. Yongquan Wu (Shanghai University).

Reference

[I] A.R. Marder. "The metallurgy of zinc-coated steel". Progress in Materials Science 45 (3) (2000),

191-271

[2] N.G Cho. Method of manufacturing a coated steel, US5571566, 1996.

[3] E.T. McDevitt. Aluminum-zinc alloy composition comprising spangle for hot-dipping steel

product, method and product obtainable thereof. EP1428898, 2007.

[4] B. Xu, D. Phelan, R. Dippenaar. "Role of silicon in solidification microstructure in hot-dipped

55wt% Al-Zn-Si coatings". Materials Science and Engineering: A 473 (1-2) (2008), 76-80

[5] K. Honda, K. Ushioda, W. Yamada. "Nucleation of the primary Al phase on TiA13 during

solidification of a hot-dip Zn-ll%Al-3%Mg-0.2%Si coating on steel sheet". Ceramic Transactions 21

(2009), 355-362

[6] B. Xu. Nucleation and growth of 55%A1-Zn alloy on steel substrate. Wollongong: University of

Wollongong, 2006.

[7] F. Garcia, A. Salinas, E. Nava. "The role of Si and Ti additions on the formation of the alloy layer

at the interface of hot-dip Al-Zn coatings on steel strips". Materials Letters 60 (6) (2006), 775-778

[8] F. Garcia, A. Salinas-Rodriguez, E. Nava. "The Role of Ti Inoculation of Al-Zn-Si Coating Alloys

on the Formation of Intermetallic Compounds by Interaction with Solid Steel". Materials Science

Forum 560 (2007), 97-102

[9] G. Wu, J. Zhang, Q. Li, K. Chou, X. Wu. "Microstructure and Thickness of 55 pet Al-Zn-1.6 pet

Si-0.2 pet RE Hot-Dip Coatings Experiment, Thermodynamic, and First-Principles Study".

Metallurgical and Materials Transactions B (2011), DOI: 10.1007/sll663-11011-19578-11662

[10] M. Guttmann. "Reactive Phase Formation at Interfaces and Diffusion Processes". Materials

Science Forum 155-156 (1994), 527-548

[II] E. Mcdevitt, Y. Morimoto, M. Meshii. "Characterization of the Fe-Al Interfacial Layer in a

Commercial Hot-dip Galvanized Coating ". 1SIJInternational 37 (8) (1997), 776-782

47

[12] K.-K. Wang, L. Chang, D. Gan, H.-P. Wang. "Heteroepitaxial growth of Fe2A15 inhibition layer in

hot-dip galvanizing of an interstitial-free steel". Thin Solid Films 518 (8) (2010), 1935-1942

[13] G Wu, J. Zhang, Y. Ren, G Li, X. Wu, Q. Li, K. Chou. "Investigation of Ti Addition Effects on

Thickness of 55%Al-Zn-1.6%Si Coating by First-Principles Calculation". Metallurgical and Materials

Transactions B (2011 ), Accepted

[14] G. Kresse, J. Furthmuller. "Efficient iterative schemes for ab initio total-energy calculations using

a plane-wave basis set". Physical Review B 54 ( 16) ( 1996), 11169-11186

[15] P.E. Blochl. "Projector augmented-wave method" Physical Review B 50 (24) (1994),

17953-17979

[16] J.P. Perdew, JA. Chevary, S.H. Vosko, K.A. Jackson, MR. Pederson, DJ. Singh, C. Fiolhais.

"Atoms, molecules, solids, and surfaces: Applications of the generalized gradient approximation for

exchange and correlation". Physical Review B 46 (11) (1992), 6671

[17] L. Stenberg, R. Sjövall, S. Lidin. "On the Compound Coordination Polyhedron in MnA16and

Fe2A15". Journal of Solid State Chemistry 124 (1) (1996), 65-68

[18] R.W. Richard, R.D. Jones, P.D. Clements, H. Clarke. "Metallurgy of continuous hot dip

aluminizing". International Materials Reviews 39 (5) (1994), 191-212

[19] K.M. Carling, E.A. Carter. "Effects of segregating elements on the adhesive strength and structure

of the a-Al203/ß-NiAl interface". Acta Materialia 55 (8) (2007), 2791-2803

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structure of the CoFeB/MgO interface from first principles ". Applied Physics Letters 89 (14) (2006),

142507-142510

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6155-6167

48

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THE ROLES OF DIFFUSION FACTORS IN ELECTROCHEMICAL CORROSION OF TiN AND CrN (CrSiCN) COATED MILD STEEL AND

STAINLESS STEEL

Feng Cai '•*, Q. Yang2, X. Huang '

1 MAE, Carleton University, 1125 Colonel By Dr., Ottawa, ON K1S 5B6, Canada 2IAR, National Research Council, 1200 Montreal Rd., Ottawa, ON K1A 0R6, Canada

Keywords: PVD coating, coating defect, corrosion, electrochemical impedance, EIS.

Abstract

Applying noble coatings on steel components is an effective solution to preventing corrosion attacks. However, through-coating defects, such as pin holes, voids and growth defects play a detrimental role in the degradation of a coating-substrate system. Through the defects, corrosive media are able to reach the metal substrate, initializing pitting corrosion and eventually resulting in coating failure. This research studies the correlation between coating defects and corrosion behavior of the TiN and CrN (CrSiCN) coated mild steel and stainless steel. Electrochemical impedance technique is used to reveal the corrosion behavior. The results revealed that in a coating-substrate system, two critical factors controlling the corrosion resistance, the effective diffusion coefficient and diffusion layer thickness, which are found to be related to coating microstructure. Denser and thicker coating structures result in lower effective diffusion coefficients and greater effective diffusion layer thickness, and demonstrate high electrochemical impedance and resistance to electrochemical corrosion.

1. Introduction

Hard coatings, especially transition metal based nitride coatings, have been extensively used to protect metal components from erosion and wear due to their excellent tribological properties [1, 2]. Also, being inert to chemicals, they are able to provide corrosion protection to less noble metals, such as steels. However, inherent microscopic defects in physical vapour deposited coatings, such as pinholes, pores and growth defects [3, 4] provide access for corrosive media to reach substrates and initiate localized corrosions at the defective sites. Such localized corrosion may result in more severe damages than that occurring on a bare metal substrate [5, 6], In Ulis study, TiN/MS (mild steel) and TiN/SS (stainless steel) coating samples were tested in 3.5 wt.% NaCl aqueous solution. The coated samples were also subject to the post-deposition treatment with polymethyl methacrylate (PMMA, [C502H8]n) to seal coating defects. PMMA has a good resistance to chemicals [7]; PMMA solution is able to fill in nano-scaled pores through in situ polymerisation of monomer MMA [8]. Thus PMMA treated coating systems are expected to have improved corrosion resistances.

' Corresponding author: Department of Mechanical and Aerospace Engineering, Carleton University, 1125 Colonel By Drive, Ottawa, Ontario, K1S 5B6, Canada, Tel: 001 (631) 828-2748; E-mail: fcai@connect.carleton.ca

49

Electrochemical impedance spectroscopic (EIS) and electrical equivalent circuit (EEC) techniques were employed to investigate the electrochemical corrosion behaviors of the TiN/MS and TiN/SS coating systems as well as the effectiveness of PMMA treatment.

2. Experimental

17-4PH stainless steel and mild steel substrates (16 mm diameter and 1.6 mm thickness) were polished to mirror finish with 0.5 urn diamond paste. TiN and CrN (CrSiCN) PVD coatings were deposited using cathodic arc (CA), plasma enhanced magnetron sputtering (PEMS), and plasma assisted electron beam (EB) deposition techniques. Some of the coated samples were subjected to post-deposition treatment with PMMA. The PMMA solution consisted of 16.6 wt.% poly(methyl methacrylate)(PMMA, (C502H8)n), 78.4 wt.% methyl methacrylate(MMA) and 5 wt.% benzoyl peroxide. After 24 hour room temperature immersion in PMMA solution, the samples were cleaned with a dry soft tissue to remove residual solution; and then cured 8 hours in ultraviolet light. Coating surface morphologies were examined using scanning electron microscopy (SEM), and the chemical compositions were measured using an energy dispersive X-ray spectroscopy (EDS). Crystal structures and grain sizes were studied using Rigaku X-ray difrractometer (XRD) with a Cu Ka radiation source; the X-ray scans were taken in the 26 range of 20° to 100° at a scan rate of2.47min. The corrosion performances of the coating samples were studied using EIS with a three-electrode cell. All the tests were conducted in 3.5 wt.% NaCl aqueous solution at room temperature. Process control and data acquisition were carried out using a Gamry potentiostat/galvanostat/ZRA system. EIS spectra were performed with an alternate current (AC) signal of 10 mV peak to peak in amplitude superimposed on the open circuit potential. The frequency range of the AC signal was set from 300 kHz to 10 mHz. The test data were interpreted through curve fitting with equivalent electrical circuit (EEC) using Gamry Echem Analyst program.

3. Results and Discussion

SEM images for the surface morphologies of the TiN coatings deposited using EB and CA techniques revealed that the surfaces of the TiN coatings were smooth with occasional particulate defects (Figure 1 a), while CA TiN (Figure 2 a) coating had a much higher defect density. This was associated with the formation of droplets in the cathodic arc process which typically generates a rougher surface [9, 10]. A typical type of defects is an embedded nodule (Figure 1 b and Figure 2 b), showing apparent gaps between the nodules and bulk coating (Figure 1 b and Figure 2 b), which can act as access channels for corrosive media to reach the substrate. In a deposition process, a high substrate bias voltage benefits atom mobility and increases surface diffusion [11, 12]. The CA TiN coatings were subjected to a high value of substrate bias voltage -200 V and a high substrate temperature of 400°C. As a result, the growth defects appeared well bonded to the coating with smaller gaps between them (Figure 2 b) compared to the EB counterpart (Figure 1 b).

50

Figure 1 SEM images of the as-deposited EB TiN coating.

Figure 2 SEM images of the as-deposited CA TiN coating.

PMMA-treated TiN coating samples were examined using SEM, and were found that PMMA filled in defect gaps without covering the bulk coating surface. SEM surface morphologies study revealed that four CrN based coating samples displayed cellular structures in addition to the presence of surface defects (embedded nodules with sizes ranging from 2 to 15 um)(Figure 3). Other defects include voids or notches along grain boundaries. Due to the coarse grain structures in the coatings CrSiCN(2) and CrSiCN(3), more voids/notches were observed. Chemical compositions of the CrN based coatings were examined using EDS, and their microstructures were studied using XRD. It was observed that coating grain size change with content of Si. With a slight addition of Si (1.3 at.%), grain size decreased from 7.6 nm (CrN coating) to 5 nm (CrSiCN(l)). Further increasing Si, grain size increased to 20.4 nm (CrSiCN(2)) and 44.5 nm (CrSiCN(3)). This trend of grain size change is in agreement with the changes in surface morphologies with content of Si, as CrSiCN(2) coating shows a coarse grain structure (Figure 3 c), and CrSiCN(3) coating shows a coarse facet structure (Figure 3 d). EIS plots measured after 24 hours of immersion are summarized in Figure 4. In the Nyquist plots as shown in Figure 4 (a), the small semicircle (arc) for the CA TiN/MS indicates a typical charge transfer controlled corrosion behavior and low diffusion impedance; the straight line

51

shown for CA TiN/SS system reveals a diffusion dominated corrosion behavior and high diffusion impedance. The different corrosion characteristics stemmed from the different corrosion mechanisms of the substrates. The corrosion on stainless steel is diffusion controlled due to the formation of a passive oxide film; while the corrosion on mild steel is in a charge transfer controlled mode because of anode dissolution. For the EB TiN/SS coating system, semicircle (arc) in the Nyquist plots revealed a charge transfer featured corrosion behavior; its diffusion impedance was between those of CA TiN/MS and CA TiN/SS coating systems. Its higher diffusion impedance comparing to CA TiN/MS was attributed to the stainless steel substrate that provided a protection with a passive oxide film; whereas its lower diffusion impedance comparing to CA TiN/SS was due to its thinner coating thickness (1.5 um), which will be discussed later. High bias voltages applied in the CA process resulted in higher surface atom diffiisivity, led to fewer through-coating voids, and eventually a lower effective diffusion coefficient. All the PMMA-treated TiN coating samples show increased diffusion impedances (Figure 4 a), which were attributed to that PMMA indeed clogged through-coating defects. In CrN based coating systems, CrN and CrSiCN(l) showed significant higher impedances than CrSiCN(2) and CrSiCN(3) (Figure 4 b). The high diffusion impedances of the CrN and CrSiCN(l) coating were attributed to their denser microstructures. In contrast, the very low diffusion impedances of the CrSiCN(2) and CrSiCN(3) coating systems were due to their larger grain sizes and porous microstructures; which will be analyzed later.

Figure 3 SEM images demonstrating morphologies of CrN based coating systems.

52

i J

Zre„, Qcm2

E o a

0

Ï

Zreal, Qcm2

(b) CrN based coating systems Figure 4 EIS (Nyquist) plots of test coating systems (at 24 hr).

To further analyze the roles of diffusion factors in the corrosion processes, EEC analyses were conducted with EIS test data. Considering the influence of coating microstructure on diffusion characteristics, Warburg diffusion and Open Finite-Length Diffusion (OFLD) EEC models were respectively applied to TiN and CrN based coating systems for infinite length diffusion and finite length diffusion processes.

53

Warburg impedance corresponds to a (semi-)infinite length diffusion process [13, 14], and is described by Warburg equation [15]:

Zw = G 14a - jo 14o> (1)

and modulus of Warburg impedance is:

[zJ^JT-alJm (2)

where, co: radial frequency of the potential perturbation; and a: Warburg coefficient which can be expressed as [15]:

RT 2F2Aj2 Coy Do CRyjDR ^

(3)

where, D0: diffusion coefficient of oxidant; DK: diffusion coefficient of reductant; A: surface area of the electrode; F: Faraday's constant, F = 96,485 Coulombs; z: number of electrons transferred; C0: bulk concentration of the diffusing oxidant; and Cfl: bulk concentration of the diffusing reductant. Warburg impedance (\Z»\) (Eq. 2) is inversely related to the diffusion coefficients of species (Do and DR) (Eq. 3). As for a finite length diffusion (OFLD) process, diffusion impedance is expressed as [16]:

Z0{m) = Zw ■ tanh[Bjj<ä] (4)

where, co: angular frequency of potential perturbation; B: characteristic diffusion parameter and is defined as [16]:

B = 8ljb~ (5)

where, D: diffusion coefficient (area/time); <5: diffusion length in general [16]; with respect to a diffusion layer, it represents the thickness of the diffusion layer (referred as to Nernst diffusion layer - NDL)[17]; in the case of a coated metal system, it can be approximately treated as the coating thickness [18]. From Eq. (4) and Eq. (5), the modulus of OFLD impedance is written as:

tanh[-==Jyffl] (6)

Eq. (4) and Eq. (6) indicate that OFLD impedance is a function of frequency, diffusion coefficient and diffusion length. At a given frequency, OFLD impedance (\Zo\) decreases due to either a high diffusion coefficient (D) or a thin diffusion layer (<5); and conversely, \Z0\ increases due to either a low diffusion coefficient (D) or a thick diffusion layer ($).

54

In a coating system, diffusion occurs through the pores in the coating. Effective diffusion coefficient was introduced to describe the through-pore diffusion, and expressed as [19]:

r

where, D: diffusion coefficient in the fluid that fills pores; e,: porosity available for transport; c: constrictivity; r: tortuosity. Eq. (7) reveals a proportional relationship between porosity in a given coating system and the corresponding effective diffusion coefficient (De) through the coating. Thus, a coating with a high porosity corresponds to a high effective diffusion coefficient (D„) (Eq. 7), and is responsible for a low Warburg diffusion coefficient (<r) (Eq. 3) and consequently low Warburg impedance \Zy\ (Eq. 1 and Eq. 2), also a low OFLD impedance (\Zo\) (Eq. 6). With regard to diffusion length or diffusion layer thickness, for a coating system with certain porosity, actual diffusion layer thickness to the substrate is related to the coating thickness [18]. A thick coating offers a thick diffusion layer to substrate, thus results in increased OFLD impedance (\Zo\) (Eq. 6).

Table 1 Summaries of coating thickness, porosity conditions, and order of magnitude in diffusion impedance (Z) for the test coating systems

Coating system

EB TiN/SS

PMMA EB TiN/SS

CA TiN/SS

PMMA CA TiN/SS

CA TiN/MS

PMMA CA TiN/MS

CrN

CrSiCN(l)

CrSiCN(2)

CrSiCN(3)

Thickness (urn)

1.5

1.5

2.8

2.8

2.8

2.8

20.2

18.9

21.1

19.6

Porosity condition

Lowered by PMMA treatment

Lowered by PMMA treatment

Lowered by PMMA treatment

Relative dense

Relative dense

Porous

porous

Z(cm2£îVs) 1.85x10- 1.06xl0+3

2.84*10+5-1.77*10+<;

6.77xl0+3-1.73*10+5

3.78xl0+ 3-1.38xl0+ 5

7.57xl0-2.87xl0+ 2

3.56xl0+2-1.6*10+4

1.51xl0+3-5.19xl0+3

5.08xl0+ 2-2.83xl0+ 3

1.62-5.09

1.03- 4.50

EEC fitting results are summarized in Table 1. With a thicker coating thickness, the CA TiN/SS sample showed about two orders of magnitude higher diffusion impedance than the EB TiN/SS sample. This was attributed to the increased electrochemical diffusion length due to thicker coating thickness; and revealed that coating thickness has a direct influence on diffusion impedance, and a sufficient coating thickness is needed to increase the electrochemical diffusion impedance for the coating system. PMMA-treated TiN coating samples showed significantly increased electrochemical diffusion impedances. This was due to the through-coating defects were clogged, thus effective diffusion coefficients were lowered. The results demonstrate the role of coating defect in governing

55

electrochemical diffusion impedance; and suggest that a low defect rate is a must for a high corrosion resistance of a coating system. In CrN based coating systems, the CrSiCN(2) and CrSiCN(3) samples showed one to two orders of magnitude lower electrochemical diffusion impedance than the CrN and CrSiCN(l) samples. These results were attributed to larger grain sizes and porous microstructures that led to higher effective diffusion coefficients through the coatings.

4. Conclusion

This work studied roles of diffusion factors, i.e., effective diffusion coefficient and diffusion layer thickness, and their correlation between the influence of coating defects and thickness on corrosion behavior of the TiN and CrN (CrSiCN) coated mild steel and stainless steel systems. With EIS test data, EEC fitting studies revealed that effective diffusion coefficient and diffusion length played controlling effects on corrosion behavior of a coating system; and in a coating/substrate system, effective diffusion coefficient is proportional to coating porosity, and diffusion layer thickness increases with coating thickness. This result provides an explanation that a denser and thicker coating structures are expected to have lower effective diffusion coefficients and greater diffusion layer thicknesses, thus offer a higher electrochemical impedance to corrosion.

References

1. N. Vershinin et al., Surf. Coat. Technol. 125 (2000) 229. 2. G. Bertrand, H. Mahdjoub and C. Meunier, Surf. Coat. Technol. 126 (2000) 199. 3. S. H. Ahn et al., Surf. Coat. Technol. 177 (2004) 638. 4. V.K.W. Grips et al., Electrochim. Acta 51 (2006) 3461. 5. H. A. Jehn, Surf. Coat. Technol. 125 (2000) 212. 6. H. Uchida and M. Yamashita, Vacuum 59 (2000) 321. 7. F. Mark et al., Encyclopaedia of Polymer Science and Technology, John Wiley & Sons, 1985. 8. J. Jang et al., Chem. Commun. (2001) 83. 9. R. Aharonov et al., Surf. Coat. Technol. 82 (1996) 334. 10. L. Hultman and J.E. Sundgren, in R.F. Bunshah (Ed.), Handbook of Hard Coatings, Noyes,

2001. 11. D.M. Mattox, Handbook of Physical Vapor Deposition (PVD) Processing, Noyes, 1998. 12. I.I. Aksenov, V.A. Belous, and V.M. Khoroshikh, XVIIth Int. Symp. Discharges and

Electrical Insulation in Vacuum, Berkeley, CA, 1996, 895. 13. C. Liu et al., Corros. Sei. 45 (2003) 1243. 14. JR. Macdonald, Impedance Spectroscopy—Emphasizing Solid Materials and System, John

Wiley & Sons, 1987. 15. A.J. Bard and L.R. Faulkner, Electrochemical Methods, Fundamentals and Applications,

Wiley, 2000. 16. F. Mansfeld et al., Electrochim. Acta 43 (1998)2933. 17. T.K. Rout, Corros. Sei. 49 (2007) 794. 18. V.K.W. Grips et al., Thin Solid Films 514 (2006) 204. 19. P. Grathwohl, Diffusion in Natural Porous Media: Contaminant Transport, Sorption /

Desorption and Dissolution Kinetics, Kluwer Academic Publishers, 1998.

56

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF ELECTROPLATING PARAMETERS ON "HER" CURRENT DENSITY IN Ni-MoS2 COMPOSITE PLATING

Ebru Saraloglu Gtller1, Ishak Karakaya1, Erkan Konca2

'Department of Metallurgical and Materials Engineering, Middle East Technical University, Ankara, Turkey

department of Metallurgical and Materials Engineering, Atilim University, Ankara, Turkey

Keywords: Electrocodeposition, Voltammogram, Current density

Abstract

Nickel composites with co-deposited insoluble, solid lubricant particles such as M0S2 have been reported to reduce friction. It is known that hydrogen evolution reaction (HER), competes with nickel deposition. The influence of the electroplating parameters and their interaction effects on the peak current density for HER were studied by fractional factorial design. The parameters and their ranges were; M0S2 concentration (0-30 g/1), temperature (30-50°C), pH (2-4) and surfactants (0-1 g/1). Electrodeposition processes were carried out from a typical Watts bath containing leveler, wetting agent and brightener by using a potentiostat. The peak currents (Ip) were extended to higher values and the peaks on linear sweep voltammograms became noticeable by increasing the scan rate from 20 mV/s to 100 mV/s over the range 0 to 2.5 V. The peak current densities (ip) for each experimental route were determined by fractional factorial design for three types of mineral processing surfactants; sodiumlignosulfonate (SLS), depramin-C (DC) and ammoniumlignosulfonate (ALS) using Minitab program [1]. Adding MoS2, decreasing temperature and increasing pH has decreasing effect on peak current density for all surfactants. ALS and DC have increasing effect whereas SLS has descending effect on peak current.

Introduction

The electrocodeposition of insoluble inert solid particles embedded in a metal matrix have been used for several industrial applications. Composite coatings based on nickel with build-in particles such as lubricants; BN [2, 3], M0S2 [4-7], polytetrafluoroethylene (PTFE) [8, 9], and graphite [10, 11] were studied to obtain decreased frictional properties. The effects of electroplating parameters such as current density, concentration of M0S2 in electrolyte, pH, temperature and rotation rate or surfactant concentration on M0S2 amount in the coating were investigated in MoS2-Ni composite coatings [5, 7]. Fractional factorial design has been used to evaluate the influence of all or some of the above parameters and also the complex interactions on corrosion resistance, deposition efficiency [12] and hydrogen evolving activity [13, 14].

Hydrogen evolution reaction (HER) which involves codeposition of hydrogen on the cathode, makes materials susceptible of hydrogen adsorption which then affects the behavior of materials in service, decreases current efficiency and leads to dull and un-uniformed surface. In nickel electroplating, the electroplating parameters influence HER and thus hydrogen permeation.

57

However, these parameters may also affect crystal axis of texture, defect presence, grain size and internal stress which can change diffusion and hydrogen trapping in nickel [15]. Texture formation of nickel deposits has been also attributed to HER in literature. Fritz et. al. claimed that increasing mean current density changes texture from (110) to (100) [16]. The relationship between texture and current density together with pH of Watts' electrolyte was in general agreement [17], but there are some differences in the values of the current densities. TEM micrograph interpretations claims that there is free mode of growth in (100) orientation due to the oblique twins and less number of dislocations whereas the characterization of (110) is vertical twins with high number of dislocations as reported by Psarrou et. al. [18]. At low current density the (110) texture was favored due to the inhibition of growth which was attributed to hydrogen adsorption that occurs at low current densities [18] whereas (100) texture was observed at the high current densities because of uninhibited growth (free growth mode) of nickel [16, 17, 19].

The relative surface energy is a measure of wetting behavior of materials one another, thus control the adherence of the electroplated materials [20]. The lowest surface energy belongs to (111) surface, then (100) and then (110) without hydrogen adsorption [21, 22]. The surface energies of nickel decrease with the adsorption of hydrogen and this decrease is more effective in (100) that lead to change in the nickel deposit texture from (111) to (100) [21, 23].

In practice, strike nickel coatings are performed at low current densities for stronger bonding [24-27]. This can be attributed to the hydrogen codeposition that favors (110) texture. The low current density application in nickel plating was also reported to assist bonding process in electronic packaging [24].

It is expected that formation of strong bonding for M0S2 particles embedded in nickel matrix may be attained by producing nickel coating that favors (110) texture. For this reason, attention was turned to HER. Therefore, the effect of the plating variables, during M0S2 particles built in nickel matrix electrocodeposition, on hydrogen evolution was studied by fractional factorial design using Minitab program [1]. Further, interaction effects were evaluated for the parameters studied using three different commonly used mineral processing surfactants.

Experimental

304 stainless steel plates (4 cm x 4 cm) were used as substrates. Substrates were placed in the electroplating cell schematically shown in Figure 1 followed by immersion into an alkaline solution and anodic acid cleaning to produce clean surface required for electrodeposition [28-30], A typical Watts bath (300 g/1 NiS04.6H20, 50 g/1 NiCl2.6H20, and 40 g/1 boric acid) was used. Nickel was coated with and without M0S2 based on the conditions determined in the statistical design given in Table I and the values were attributed to the low and high levels by interpreting the studies in literature [4, 5, 7]. To identify the effects of 5 electroplating variables on hydrogen evolution profile with the reasonable number of experiments, 25"1 fractional factorial design was used in the statistical design. The variables were determined as A: M0S2 concentration, B: Temperature, C: pH, D: Surfactant, E: Coating thickness and their fixed values are listed in Table I. The peak current (Ip) and the peak voltage (Vp) values in the process were chosen as response values in the design. Potentiometric measurements were carried out and

58

"echem analyst" program was used to determine the Ip and Vp values. The potential difference between the working electrode (cathode) and the nickel anode was swept linearly from 0 to 2.5 V at scan rates of 20 mV/s and 100 mV/s.

Figure 1. Schematic view of the electrodeposition set-up

Table I, Factors and Levels for Fractional Factorial Design (FFD)

Factor

A M0S2 cone (g/1) B Temp(C) C pH D Surfactant (g/1) E Coating Thickness (um)

Level -1 0 30 2 0 0

+1 30 50 4 1 40

Results and Discussion:

25"' fractional factorial design using Minitab program [1] gives the experimental route in Table II. Voltammetric response associated with HER during depositing Ni and Ni-MoS2 composite from Watts' bath onto the stainless steel electrode was focused. At the beginning, potential was swept at two different scan rates; 20 mV/s and 100 mV/s. The effect of scan rate on the voltammogram is illustrated in Figure 2 for experiment 2 conditions of Table II. It is apparent that total current at extremum increases with increasing scan rate as expected [31]. 100 mV/s was chosen as the scan rate in the following experiments since Ip's were well determined and predominant as potential scan rate increased. From this figure it can be seen that the current is almost zero until the onset of reduction of nickel and/or hydrogen ions at the cathode. Then it continues to increase and forms a maximum. The irregularity of the voltammogram after maximum indicates the decay of one of the primary electrode reactions. To test the decay of HER, current efficiencies of nickel plating were determined at 3 current densities selected from the voltammogram. One of them was the peak value. The other two were before and after die peak value. Under the conditions of experiment 15, for 2 and 8 hours of depositions, the average percent current efficiencies were determined as 90, nearly 100 and 68 for current densities before, after and at the peak respectively. The decrease in the current

59

efficiency at the peak position can be attributed to HER since cathodic current efficiency of deposition in Watts' bath decreases due to the hydrogen evolution nearly at a current density of 0.65 A/dm2 [32]. The peak current densities, shown in Table III, are comparable to above value for experiments performed in this study.

Table II. Experimental conditions for 2 "' Fractional Faci values for the variab!

Experiment No

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15 16

A (MoS2) c

-1 -1

-1

-1

-1 -1 -1

- ]

B (Temperature)

-1

-1

-1 -1

-1

-1

-1 -1

C (pH)

-1

-1 -1

-1 -1 -1 - ]

-1

D (Surfactant)

-1 -1

-1

-1

-1 -1

-1 -1

E (Thickness)

-1 -1 -1 -1 -1 -1 -1 -1

orial Design: es

1 = low values l=high

Figure 2. The effect of scan rate on the voltammogram for experiment 2 conditions of Table 2.

60

The typical volammograms and the positions of Ip values for the first four experiments were indicated in Figure 3. The values of the peak current, Ip, the peak current density, ip, and the peak voltage, Vp, together with surfactant designations are given in Table III for all experiments.

Figure 3. The typical volammograms at 100 mV/s scan rate and the positions of Ip values, in mA, for the first four experiments

The results were subjected to factorial design analysis to obtain interaction effects of surfactants as illustrated in Figure 4. Four parameters were used, since coating thickness has no effect on current density. The parameters that have reducing effect on the current density yielded negative coefficients in the regression analysis (Equations 1, 2 and 3) and the magnitudes state the level of their effects. The corresponding surfactant designations of each equation are given inside the parenthesis next to the equations.

ip (A/dm2) = 0,465 - 0,0259 A + 0,0112 B - 0,149 C - 0,0752 D (SLS) ( 1 ) ip (A/dm2) = 0,557 - 0,0486 A + 0,0814 B - 0,204 C + 0,0167 D (DC) (2) ip (A/dm2) = 0,675-0,113 A+ 0,095 B-0,186 C +0,135 D (ALS) (3)

The variables; A, B, C and D given in above equations are presented with their units given in Table I. It can be concluded that parameters A and C have decreasing but B has increasing effects on the peak current for all surfactants. If ALS and DC are used increasing effect whereas if SLS is used as surfactant descending effect on peak currents are observed.

61

Table III. I„ (cathode 10 cm ) ,VP and ip values for experiments from l to 16 Experiment 1 2

3

4

5

6

7

8

9

10 11

12

13

14 15

16

Surfactant

SLS DC ALS

SLS DC ALS SLS DC ALS

SLS DC ALS

SLS DC ALS SLS DC ALS

SLS DC ALS

In (A) 0.07040 0.03950 0.04686 0.05772 0.06210 0.07079 no peak 0.03944 0.05628 0.02887 0.02672 no peak no peak 0.03541 0.01056 0.01412 0.01960 0.05273 0.04745 0.03090 0.07180 0.04009 0.02182 0.02053 0.03340 0.04153 0.04187 0.03243 0.01989 0.02032 0.02084 0.03631

v.m 2.323 1.861 2.027 1.926 2.059 2.160 no peak 1.879 1.728 2.027 1.865 no peak no peak 2.027 1.915 2.142 1.572 1.686 1.716 1.609 2.023 1.866 1.498 1.535 1.629 1.676 1.719 1.649 1.595 1.860 1.592 1.706

UA/dm2) 0.704 0.395 0.469 0.577 0.621 0.708 no peak 0.394 0.563 0.289 0.267 no peak no peak 0.354 0.106 0.141 0.196 0.527 0.474 0.309 0.718 0.4 0.218 0.205 0.334 0.415 0.419 0.324 0.199 0.203 0.208 0.363

Conclusions The peak currents (Ip) during hydrogen evolution reaction (HER) in Watts nickel plating bath, were determined from linear sweep voltammograms by using 100 mV/s scan rate over the range 0 to 2.5 V. According to the interaction plots obtained from Minitab program [1], it was deduced that adding M0S2, decreasing acidity and temperature lead to a decrease in ip for all three surfactants. Moreover, addition of SLS decreased, but ALS and DC increased ip.

62

Figure 4. Interaction plot for peak currents; the columns are showing; (a) A-B interaction (b) A-C and B-C interactions and A-D, B-D and C-D interactions are for (c) Surfactant SLS (d)

Surfactant DC and (e) Surfactant ALS. The dashed and full lines in this figure represent high and low levels respectively. The ip scales of vertical axes are not shown for simplicity.

References 1. Minitabl5,"Statistical Software" 2. E. Pompei, et al.,"Electrodeposition of nickel-BN composite coatings,"Electrochimica Ada, 54 (2009), 2571-2574. 3. N.K. Shrestha, et al./'Composite coatings of nickel and ceramic particles prepared in two steps,"Surface and Coalings Technology, 140(2001), 175-181. 4. L.M. Wang,"Effect of surfactant BAS on MoS2 codeposition behaviour,"Journal of Applied Electrochemistry, 38 (2008), 245-249. 5. Y.C. Chang, Y.Y. Chang, and C.I. Lin,"Process aspects of the electrolytic codeposition of molybdenum disulfide with nickel,"Electrochimica Ada, 43 (1998), 315-324. 6. M.F. Cardinal, et al.."Characterization and frictional behavior of nanostructured Ni-W-MoS2 composite coatings "Surface and Coatings Technology, 204 (2009), 85-90. 7. S.L. Kuo,"The Influence of Process Parameters on the MoS2 Content of Ni-MoS2 Composite Coating by the Robust Design Method, "Journal of the Chinese Institute of Engineers, 27 (2004), 243-251. 8. Y. Kunugi, et al.."Electrolysis using composite-plated electrodes: Part III. Electroorganic reactions on a hydrophobic Ni/PTFE composite-plated nickel electrode, "Journal of Electroanalytical Chemistry and Interfacial Electrochemistry, 313 (1991), 215-225. 9. P. Berçot, E. Pefla-Mufioz, and J. Pagetti,"Electrolytic composite Ni-PTFE coatings: an adaptation of Guglielmi's model for the phenomena of incorporation,"Skrface and Coatings Technology, 157 (2002), 282-289. 10. F. Plumier, et al.."Electrolytic co-deposition of a nickel/fluorographite composite layer on polycrystalline copper,''Applied Surface Science, 212-213 (2003), 271-278. 11. P. Dabo, H. Ménard, and L. Brossard,"Electrochemical characterization of graphite composite-coated electrodes for hydrogen evolution reaction,"'Internationaljournal of Hydrogen Energy, 22 (1997), 763-770.

63

12. R. Santana, et al.."Studies on electrodeposition and corrosion behaviour of a Ni-W-Co amorphous alloy"Journal of Materials Science, 42 (2007), 9137-9144. 13. C.C. Hu and A. Bai,"Optimization of Hydrogen evolving activity on Nickel-Phosphorus Deposits using Experimental Strategies, "Journal of Applied Electrochemistry, 31 (2001), 565-572. 14. C.C. Hu and C.Y. Weng,"Hydrogen evolving activity on nickel-molybdenum deposits using experimental strategies,"Jo«rna/ of Applied Electrochemistry, 30 (2000), 499-506. 15. L. Mirkova, et al.,"Hydrogen coevolution and permeation in nickel electroplating,"Jowna/ of Applied Electrochemistry, 33 (2003), 93-100. 16. T. Fritz, W. Mokwa, and U. Schnakenberg,"Material characterisation of electroplated nickel structures for microsystem technology,"£/ecf/"ocW/w(ca Ada, 47 (2001), 55-60. 17. T. Fritz, et al.."Characterization of electroplated nickel/'M/crasys/e/M Technologies, 9 (2002), 87-91. 18. S. Psarrou, P. Gyftou, and N. Spyrellis,"Electron Microscopy Study of Nickel and Nickel Composite Electrocoatings,"Mcroc/i;/wica>4cto, 136 (2001), 159-163. 19. R. A.K.N,"Preferred orientations in nickel electro-deposits: I. The mechanism of development of textures in nickel electro-deposits,"./owrta/ of Electroanalytical Chemistry (1959), 6 (1963), 141-152. 20. H. Brown and B.B. Knapf. "Modern Electroplating," Electroless and Electrodeposition of Silver, ed. F.A. Lowenheim (Lowenheim: Wiley-Interscience, 1974), 137. 21. Y.Y. Huang, Y.C. Zhou, and Y. Pan,"Effects of hydrogen adsorption on the surface-energy anisotropy of nickel,"Physica B: Condensed Matter, 405 (2010), 1335-1338. 22. D.Y. Li and J.A. Szpunar,"A Monte Carlo simulation approach to the texture formation during electrodeposition—I. The simulation model,"Electrochimica Ada, 42 (1997), 37-45. 23. H. Li, F. Czerwinski, and J.A. Szpunar,"Monte-Carlo simulation of texture and microstructure development in nanocrystalhne electrodeposhs,"Nanostructured Materials, 9 (1997), 673-676. 24. L.W. Pan, L. Lin, and J. Ni,"A flip-chip LIGA assembly technique via electroplating,"Microsystem Technologies, 7 (2001), 40-43. 25. J. Youngcheol, D. Kiet, and K. Chang-Jin,"Fabrication of monolithic microchannels for IC chip cooling,"M(cro Electro Mechanical Systems. MEMS '95, Proceedings. IEEE, (1995), 362-367. 26. L.W. Pan, et al.,"Flip chip electrical interconnection by selective electroplating and bonding,"Microsystem Technologies, 10 (2003), 7-10. 27. S. Gao and A.S. Holmes,"Thermosonic Flip Chip Interconnection Using Electroplated Copper Column Arrays,"'Advanced Packaging, IEEE Transactions on, 29 (2006), 725-734. 28. M. Handbook,vol. 5 (United States of America: ASM International ( 1994), 13. 29. R.N. Gay and W.K. Raymond (1988). 30. Astm,"Standard Practice for Preparation of and Electroplating on Stainless Steel,"B 254 -92 (2004). 31. P.M.S. Monk. "Fundamentals of Electroanalytical Chemistry," Linear Sweep and Cyclic Voltammetry at Solid Electrodes, ed. England: John Wiley and Sons Ltd. (2001). 32. M. Ibrahim,"Black nickel electrodeposition from a modified Watts bath,''Journal of Applied Electrochemistry, 36 (2006), 295-301.

64

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

PRODUCTION OF CERAMIC LAYERS ON ALUMINUM ALLOYS BY PLASMA ELECTROLYTIC OXIDATION IN ALKALINE SILICATE

ELECTROLYTES

Alex Lugovskoy, Aleksey Kossenko, Barbara Kazanski, Michael Zinigrad

Ariel University Center of Samaria; Scientific Park; Ariel 40700, Israel

Keywords: Plasma electrolytic oxidation, corrosion protection, oxide layer, aluminum alloys

Abstract

Plasma electrolytic oxidation (PEO) is a technology allowing obtaining hard wear- and corrosion-resistant in the form of thick and highly-adhesive oxide layer on aluminum surfaces. The process can be performed in several types of electrolyte, of which the alkaline silicate electrolytes were employed in this study. Silicate electrolytes passivate aluminum alloys, favor micro-arc discharges and contain silicate ions, which modify the technological properties of the coating. The influence of the silicate index of sodium silicates and of their concentration in the electrolyte on the composition, structure and properties of the oxide layer was studied. Electrolyte properties, electrochemical process parameters and the properties of the resulting coating were studied and compared. Optimal electrolyte compositions for the obtaining of hard and corrosion resistant ceramic layer were found. A plausible mechanism of the process was proposed.

Introduction

Plasma electrolytic oxidation (PEO, also known as Micro-Arc Oxidation) can be roughly considered a kind of anodizing for valve metals, primarily, aluminum, magnesium, titanium and their alloys. However, there are some features allowing speaking of the PEO as of a distinct technique for the production of oxide layers on metal surfaces. While anodizing is normally performed by low-voltage DC current (AC is possible for some special cases), very high AC voltage is mandatory for the PEO (high DC voltage can be applied for some special cases). Sparking on or in the vicinity of the treated surface is an absolutely undesirable phenomenon for the anodizing; for the PEO intensive sparking is an indication of a normal process. PEO produces much thicker (as compared to anodizing) oxide layers, which are typically less susceptible to fatigue. In spite of relatively high porosity, PEO layers effectively protect the base metal against corrosion due to the fact that the pores are always capped with a barrier oxide layer and are therefore non-penetrable by corrosion media. The improvement of corrosion stability of PEO treated metals as compared to bare metals has been reported for aluminum [1, 2] and magnesium alloys [3, 4],

The main peculiarities governing plasma electrolytic processing were outlined by Yerokhin et al. [5], which proposed the mechanisms of very complex phenomena in the vicinity of the metal-electrolyte interface during plasma discharges. In spite of considerable lacunae still existing in the detailed understanding of the processes and the factors affecting the resulting coating, it is accepted by most authors that in the course of each AC period several principal step

65

occur: (1) a barrier oxide layer is formed on the boundary between the metal and the electrolyte during the initial anodic semi-period; (2) the potential difference on the two sides of the dielectric oxide layer grows as the anodic semi-period is advancing until (3) the dielectric breakdown takes place. The breakdowns through the oxide layer are accompanied by sparks so that the process actually occurs in the mode of micro-arc discharges. Fresh portions of the electrolyte are injected to the bare metal surface during the breakdowns and the process continues as long as the voltage is sufficient for new breakdowns which perforate the growing oxide layer. Metal and oxide relaxation and partial reduction of oxidized species occur in the course of cathodic semi-periods. Gas micro-phase formation (nucleation) and annihilation (cavitation) processes apparently contribute to the oxide layer production, but these have been scarcely studied because of evident experimental difficulties.

Various versions of the PEO differ from each other in the profile of the applied electricity and the composition of electrolyte. Industrial sine 50-60 Hz alternating high (e.g. ±400V nominally, +400-600V to -200V in practice, due to partial rectifying effect of the valve metal oxide) voltage is most frequently used. The produced oxide layer consists of two sub-layers: an outer brittle sub-layer typically having the hardness of 500-1000 HV and the porosity of 15% and the inner functional sub-layer whose typical hardness is 900-2000 HV and typical porosity is 2-15%. The outer brittle sub-layer can be easily removed by polishing and the inner harder sub-layer can be finished to smooth marble surface.

PEO layers can be produced in several types of electrolytes [5], whose components can dissolve the base metal (acidic and alkaline electrolytes), passivate it (phosphate and polymer electrolytes) or interact with it in more complex and less understood ways (fluoride electrolytes).

We here report the results of a study of PEO of aluminum A5052 alloy in alkaline-silicate electrolytes prepared from several silicates having different SiÛ2: Na2Û ratios (silicate indexes).

Experimental

Small rectangular (3 x 15 x 30 mm) specimens of aluminum A5052 alloy (Al as the base and approximately 2.5% of Mg) were cut, polished with #1200 grit SiC abrasive paper and rinsed in tap water prior to be PEO processed. The oxidation was performed in AC mode by the industrial 50 Hz sine voltage (± 400 V, nominally) at the end current density 6.6 ± 0.2 A / dm2

for 30-60 minutes on a home-made 40 kVA PEO station with a water-cooled bath made of stainless steel, which served as the counter electrode. Potassium hydroxide KOH (Finkelman Chemicals, technical grade), sodium silicate Na2Si03 (pentahydrate, Spectrum, practical grade), and water glass Na20-3Si02 (Spectrum, practical grade) having the silicate indexes n = 1 and n=3, respectively were used for the preparation of the electrolytes.

Conductivities and pH of the electrolytes were measured by a YK-2005WA pH/CD meter, the thickness of oxide layers was measured by a micrometer, coating thickness gauge CM-8825 and by SEM. The surface morphology, structure and composition were inspected on SEM JEOL JSM6510LV equipped with an NSS7 EDS analyzer (Correction Method Proza- Phi-Pho-Z was used for the quantitative analysis). Cross-section samples prepared according to standard metallographic protocols were used for SEM, EDS and XRD. X-ray Diffractometer (XRD) Panalytical X'Pert Pro with Cu Ka radiation (X.=0.154 nm) was used with the full pattern identification made by X'Pert HighScore Plus software package, version 2.2e (2.2.5) by PANalytical B.V. Materials identification and analysis made by the PDF-2 Release 2009 (Powder Diffraction File). Phase analysis identification made by XRD, 40kV, 40mA. The XRD

66

patterns were recorded in the GIXD geometry at a=l°and 5° in the range of 20-80° (step size 0.05° and time per step 2s).

Autolabl2 Potentiostat with a standard corrosion cell with an Ag|AgCl reference electrode was used for corrosion tests.

Results and discussion

Each electrolyte contained 1 gr/L (17.9 mmol/L) KOH and various silicates as specified in Table 1. Conductivities of the electrolytes were at least 4-5 mS/m and all the electrolytes had pH= 11-13 (see Table 1).

Table 1. Typical electrolyte parameters (pH / Conductivity, mS/m) Na20nSi02 5g/L 10g/L 15g/L

Na2Si03 (n = 1)

Na2Si307 (n = 3)

12.68/10.27

11.08/4.53

12.74/15.5

11.18/5.47

12.80/22.7

11.24/6.52

As seen from Table 1, both the basicity and the conductivity are strongly affected by the silicate index, which is not surprising because the molar fraction of sodium oxide is 0.5 for Na2Si03 and only 0.25 for Na2Si307. The values of pH of the electrolytes only weakly depend on the concentration of a given silicate, while their conductivities are roughly proportional to the concentration of Na2Si03 or Na2Si3C>7. As one could expect, better conductivities of the "n=l electrolytes" must facilitate the PEO process. Indeed, lower current densities are needed for the plasma process initiation when n = 1 (Figure 1). As seen from Figure 1, not only the initial current densities, but also the process voltages are higher for "n = 3 electrolytes".

Figure 1. Typical amplitudes of voltages (left) and current densities (right) for PEO of n=l and n=3 electrolytes.

Oxide layers produced after 30-60 min of PEO have porous morphology with blind "crater-like" pores, which are the results of plasma discharges through the oxide (Figure 2). The morphology and elemental composition of a pore obtained by the EDS are presented in Figure 3. As seen in

67

Figure 3, the interior of a pore contains much less silicon and much more aluminum than the exterior close to the surface, because the only source of silicon is the electrolyte.

Figure 2. A SEM image (Secondary electron detector) of the morphology of an oxide layer produced by PEO.

Figure 3. Morphology and elemental composition of a pore at two characteristic points: point 1 (ptl) on the surface of the specimen; point 2 (pt2) inside the pore.

68

Figure 4. The structure of oxide layers on cross-sections of specimens obtained by PEO in 0.05 mol/L n=l (left) and n=3 (right) electrolytes: (1) non-oxidized base alloy, (2) internal denser oxidized sublayer, (3) external loose sublayer, (4) resin wrapping. Back-Scattered Electron SEM image, x 1,000.

The oxide layers, formed after 30 minutes of PEO, are 20-60 um thick and consist of two clearly pronounced sublayers: a denser internal sublayer and a loose porous external sublayer (Figure 4). Both sublayers are considerably thicker for the "n=3 electrolytes", but they contain 2-5 times as much silicon as for the "n=l electrolytes". XRD phase analysis evidences that the oxide layer contains 60-70% of y-alumina, 20-30% of n-alumina and about 1% of quartz for the "n=l electrolytes." For "n=3 electrolytes" the oxide layer consists mainly of mullite 3Al203-2Si02

and varied (for different concentrations of the n=3 silicate in the electrolyte) amounts of amorphous silica, quartz and various types of alumina.

Results of hardness measurements of oxide layers are presented in Table 2. As seen from

Table 2. Microhardness of oxidized sublayers produced by PEO in electrolytes containing different concentrations of "n=l" and "n=3" silicates. (Measured on cross-sectioned specimens

perpendicularly to the section planes.) Silicate

concentration in electrolyte, mol-L"1

0.013

0.019

0.025

0.050

0.063

0.075

0.100

0.125

0.150

Microhardness, Vickers (HV) n=l electrolyte, external sublayer

840

1130

890

900

920

870

900

850

n=l electrolyte, internal sublayer

1100

1380

1050

1120

1630

1050

1060

990

n=3 electrolyte, external sublayer 1480

770

1280

710

700

n=3 electrolyte, internal sublayer 1380

1060

1570

980

910

69

Table 2, the mean microhardness for the "n=l electrolytes" is HV910 for the external and HV1200 for the internal sublayer; for "n=3 electrolytes" the corresponding values are HV990 and HV1200, respectively. No clear dependence of the microhardness on electrolyte concentration is observed.

Corrosion tests were made after a specimen was masked by resin except for a square window having the area of 1 cm2 on the oxidized surface. Thus prepared specimen was held for 1 hour in 1% NaCl and then its voltammetric curve was measured using Linear Sweep Voltammetry. Voltammetric coordinates were then replaced by Tafel coordinates (E vs. logi) and corrosion current densities and recalculated to the standard hydrogen electrode potentials were determined from Tafel slopes (Figure 5).

Figure 5. A sample of a corrosion voltammetric curve in Tafel coordinates for the determination of corrosion potential (Econ0 and corrosion current density ( w ) .

The results of thus measured corrosion characteristics of "bare" A15052 alloy and different PEO oxidized specimens are given in Table 3.

Table 3. Corrosion electrolytes.

Silicate concentration in

electrolyte, mol- L"1

"bare" AI5052

0.013

0.025

0.050

0.075

0.100

0.150

current densities

n=l electrolyte, Ecorr,Vvs.SHE

-1.126

-0.497

-0.796

-0.942

-0.995

and potentials of A15052 alloy oxidized in different

n=l electrolyte, icorr, A/Cm2 X 106

15.99

3.60

4.30

1.93

3.77

n=3 electrolyte, Ecorr, V vs. SHE

-0.525

-0.815

-0.998

-0.972

n=3 electrolyte, icon-, A/cm2 x 106

0.08

2.66

0.98

1.68

70

As follows from Table 3, corrosion current densities measured on oxidized samples are at least 3-4 times lower than for the untreated alloy. Corrosion potentials for all the oxidized samples are considerably more positive than for the untreated alloy, which evidences the increase of anodic stability in the test solution. These results are better than those obtained for anodizing [6] and similar to those obtained for much more expensive protection methods. All the measurements evidence that specimens treated in "n=3 electrolytes" have better corrosion protection than those treated in "n=l electrolytes." The latter is somewhat surprising, because mullite, which is the main constituent of PEO layers produced in "n=3 electrolytes", is considered a relatively loose coating producing only poor corrosion protection.

Conclusions Plasma Electrolytic Oxidation in alkaline silicate electrolytes containing 0.013-0.150

mol-L"1 of sodium silicate having silicate index n=l and n=3 was performed on A15052 aluminum alloy. For all the electrolytes 20-90 urn thick technological layer was obtained; composition, structure and properties of the oxidized layer were studied. For each sample, the oxidized layer consists of a denser internal and looser external sublayer. While for "n=l electrolytes" the oxidized layer is mainly formed by several kinds of alumina, the principal constituent of the oxidized layer for "n'3 electrolytes" is mullite.

Measurements of microhardness evidenced that it is apparently not influenced by the kind of silicate (n=l or n=3) and by its concentration in the electrolyte.

Electrolytes with silicate index n=3 ensure better corrosion protection that those with n=l. Corrosion protection parameters are significantly better for all PEO oxidized samples than for the untreated A15052 alloy.

References [1] R.C. Barik, J.A. Wharton,T, R.J.K. Wood, K.R. Stokes, R.L. Jones, "Corrosion, erosion and

erosion-corrosion performance of plasma electrolytic oxidation (PEO) deposited A1203 coatings," Surface & Coatings Technology, 199 (2005), 158- 167.

[2] Wenbin Xue, Xiuling Shi, Ming Hua, Yongliang Li, "Preparation of anti-corrosion films by microarc oxidation on an Al-Si alloy," Applied Surface Science, 253 (2007), 6118-6124.

[3] Hongfei Guo, Maozhong An, Shen Xu, Huibin Huo, "Microarc oxidation of corrosion resistant ceramic coating on a magnesium alloy," Materials Letters, 60 (2006), 1538-1541.

[4] Fei Chen, Hai Zhou, Bin Yao, Zhen Qin, Qingfeng Zhang, "Corrosion resistance property of the ceramic coating obtained through microarc oxidation on the AZ31 magnesium alloy surfaces," Surface & Coatings Technology, 201 (2007), 4905^*908.

[5] A.L. Yerokhin, X. Nie, A. Leyland, A. Matthews, S.J. Dowey, "Plasma electrolysis for surface engineering," Surface and Coatings Technology, 122 (1999), 73-93.

[6] C. V. Zabielski, M. Levy, "Study of Type II and Type III Anodized Al 5052-0 in Aqueous DS2 Solutions," U.S. Army Research Laboratory Environmental Effects, 1992.

[7] G. da Silva, M. Ueda, J. O. Rossi, L. L. G. da Silva, M. M. da Silva, "Corrosion Resistance and Hardness of A15052 Alloy Treated by Plasma Immersion Ion Implantation," 17° CBECIMat - Congresso Brasileiro de Engenharia e Ciência dos Materials, 15-19 November 2006, Foz do Iguaçu, Brasil.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ABRASIVE WEAR PROPERTIES OF PLASMA SPRAYED Y-PSZ COATED 6063-T6 ALUMINUM ALLOY

Eray Erzi, Selim Yildirim, Suat Yilmaz

Metallurgical and Materials Engineering Department, Istanbul University, 34320, Avcilar, Istanbul, Turkey

Keywords: Plasma spraying, Y-PSZ coating, 6063 aluminum alloy, abrasive wear.

Abstract

Thermal spray methods and especially plasma spray method have an increasing importance and being widely used by industries. This method depends on the purpose of spraying powder based coating materials to a surface. Light alloys like aluminum alloys have some disadvantages like low mechanical strength and wear resistance Aim of this study is investigating the wear resistance behavior of coated and uncoated T6 heat treated 6063 Aluminum alloys. Coatings were made by plasma spray method and %8 Y-PSZ was used as coating material. Wear tests were done by pin-on-disc method with three different load parameters and results were discussed comparatively.

Introduction

Light metals such as aluminum and its alloys are extensively used today due to weight and energy saving [1], Because of their ductile and low tribological properties aluminum tend to show low wear resistance. Numerous researchers are studying to improve and vary its usage areas however to use aluminum such as moving machine parts, aluminum has to show better mechanical strength and resist higher temperatures. One of these studies is coating aluminum alloys surface with hard ceramic coatings [2, 3]. Soft material with a harder coating performs better strength so that one of the materials to coat aluminum substrate is ceramics. With harder coating, soft materials tend to wear less as long as the coating surface exists. Thus, with the soft materials design flexibility complex parts can be produced cheaper [4-7], In this study, wt.%8 Y2O3 stabilized zirconia powders (Y-PSZ) coated on to T6 heat treated aluminum alloys surface with plasma spraying method and coated samples wear resistances were compared with uncoated samples [8].

Experimental

The composition of the samples which were used in the experiment is 6063 aluminum alloy consist of %0.40 Mg, 0.40 Si and %0.10 Cu approximately and it can be age hardened. Used samples were produced from 30mm rod which was produced with injection press. These rods then exposed to T6 aging process. Rods first heated up to 480 C then heat treated at this temperature for 20 minutes and then water quenched to room temperature. Then rods heated to 180°C and heat treated for 7.5 hours for aging. Heat treatments were done with Lenton UAF 15/10 model laboratory furnace. Aim of this heat treatment process is to gain of better machinability.

Y-PSZ ceramic powders (Y2O3) which were used as coating material were Powder Alloy Corp. product. 6063-T6 heat treated samples first surface roughened for better grip in the coating stage. Later a lining material was coated on the 6063-T6 aluminum base. The purpose of this lining is to

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bind the metal base and the ceramic coating. The lining material was coated by Sulzer-Metco 3B model plasma spray gun and the lining material was Metco Amdry 956 product and has an average size of ~45um. After lining; Powder Alloy Corp. product ~50um grain sized Yttria stabilized zirconia powders were coated with the same plasma spray gun (Figure 1).

Figure 1. Uncoated (a) and Y-PSZ coated (b) wear test samples.

Materials and Method

Abrasion tests were carried out with pin on disk method and calculated as weight loss. The weight loss was measured with ±0,0001 precision. Weights of the dry samples were measured with precision scales. By using weight loss and density of the samples volumetric losses were calculated.

AV = AG/p (1)

AV: Volumetric loss (cm3), AG: weight loss (g), p: density (g/cm3). Friction coefficient values are also calculated by wear test device [8]. The wear test device is Jinan MMW-1A model (China). Pure alumina pins are used as abrasive and samples were prepared as discs.

Results and discussion

Uncoated and Y-PSZ coated 6063 samples were both tested with pin on disk wear tests. These test were carried out with three different loads. By doing so while coated and uncoated samples were tested also the effect of the load to the wear resistance were investigated. Three different load values are used. These are ION, 20N and 40N. Wear test durations were 2.000 cycle and speed of the cycle is 60 rpm. 2.000 cycle is approx. 14,46 m. Uncoated and coated samples weight loss values for different loadings are given in Figure 2, and average values of friction coefficients are given in Figures 3 and 4.

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Figure 2. Uncoated and Y-PSZ coated 6063-T6 aluminum alloy samples abrasive wear behavior in different loadings (ION, 20N and 40 N)

Figure 3. Friction coefficient graph and average value of uncoated 6063-T6 aluminum alloy.

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Figure 4. Friction coefficient graph and average value of Y-PSZ coated 6063-T6 aluminum alloy.

It can be seen from the Figure 2 that uncoated and coated 6063-T6 aluminum alloys wear rate were increased with increasing load values. In all experiments coated 6063-T6 alloys wear resistance were better than all uncoated aluminum alloys. It can be see from the results that the coating has improved the wear resistance of 6063-T6 aluminum alloys. In the planned test conditions coated samples were performed almost 10 times better wear resistance than uncoated samples. With increasing loading both uncoated and coated samples weight loss was increased. Average values of kinetic friction coefficients in all experiment are given in Figures 3 and 4. Accordingly, the friction coefficients were determined for 10 N loading and calculated as 0,135 for uncoated and 0,148 for Y-PSZ coated 6063-T6 aluminum alloys. Friction coefficient results for 20 N and 40 N loadings similar values were calculated.

Conclusion

For ION, 20N and 40N loading wear test conditions wt. %8 Y-PSZ coating by plasma spray process, highly increases the 6063-T6 aluminum alloys wear resistance. Friction coefficient values have increased with Y-PSZ coating but this difference is very small.

Acknowledgement

This work was supported by Research Fund of Istanbul University (Project Nr. : 515/05052006).

References

1. R. Gadow, D. Scherer, "Combined coatings on light metal alloys for applications under dry sliding conditions" 10lh International Metallurgy and Materials Congress, Istanbul, Turkey, (24-28 May 2000), 861-868.

2. L. Wang, X. Nie, "Silicon effects on formation of EPO oxide coatings on aluminum alloys", Thin Solid Films, 494 (2006), 211-218.

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3. S.G. Xin, "Composition and thermal properties of the coating containing mullite and alumina," Materials Chemistry and Physics, 97 (1)(2006), 132-136.

4. K.. Kato, "Wear in relation to friction a review," Wear, 241(2000), 151-157. 5. R.C Tucker et al., "The use of thermal spray coatings for protection against wear and

corrosion," Protective Coatings and Thin Films, (1997), 387-399 6. A.P. Josep, "HVOF thermal sprayed coatings on aluminum alloys and aluminum matrix

composites," Surface and Coatings Technology, (2005), 52-56 7. G. Barbezat, "Advanced thermal spray technology and coating for lightweight engine

blocks for the automotive industry," Surface and Coatings Technology, 200 (2005), 1990-1993

8. E. Erzi, "Investigation of mechanical properties of thermally sprayed partially stabilized zirconia coated light metal alloys" (PhD thesis, Istanbul University,2011).

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THE ELECTROCHEMICAL BEHAVIOR OF SURGICAL GRADE 316L SS WITH AND WITHOUT HA COATINGS IN

SIMULATED BODY FLUD3

Tejinder Pal Singh', Harjinder singh2, Hazoor Singh3, Harpreet Singh Saheet4

'Gulzar Institute of Engineering & Technology, Ludhiana, Punjab, India Government Medical College and Rajindra Hospital, Patiala, Punjab, India

3Yadavindra College of Engineering, Bathinda, Punjab, India 4School of Mechanical, Materials and Energy Engineering, Indian Institute of Technology Ropar,

Roopnagar-140001, Punjab, India, Tel.: +91-1881-242177; fax: +91-1881- 223395. E-mail address: harpreetsingh@iitrpr.ac.in

Keywords: 316L SS, Corrosion, Thermal spraying, HA coating, Simulated body fluid, MTT assay

Abstract

ASTM recommended surgical grade 316L SS stainless steel is one of the most widely used material in orthopedic implants. It has been reported that this steel is often susceptible to pitting corrosion. The main aim of this study is to evaluate the corrosion behavior of uncoated and hydroxyapatite (HA) coated 316L SS in simulated body fluid conditions. HA coating was deposited using a thermal spraying technique. The coatings were characterized by XRD, SEM/EDS and electrochemical techniques. The corrosion resistance of the steel was found to increase after the deposition of the HA coating. The cell culture studies show that HA coated 316L SS specimens appeared more bioactive than uncoated 316L SS specimens.

Introduction

Surgical implants are usually made of metallic materials, such as austenitic stainless steel, cobalt-chromium alloys, and titanium and its alloys. Amongst all these materials, surgical grade 316L SS is the most popular material due to its relative low cost, increased mechanical strength, ease of fabrication and reasonable corrosion resistance. 316L SS is widely used as bone plates, screws and nails etc. in orthopedic applications [1-2]. Corrosion of 316L SS has been a subject of extensive research [3-5]. Oxide film on the stainless steel consists of inner chromium oxide layer and outer nickel oxide layer. However, when these materials are used in highly aggressive human body environment, their degradation takes place and corrosion products of iron, chromium nickel and molybdenum release in the body. It is reported that nickel and chromium can cause allergic problems due to their toxic effects [3], Sivakumar and Rajeswari [4] reported the failure of 316L SS intramedullary nail in the right femur of a man, due to pitting-induced stress corrosion cracking. Surface modification of metallic implants with a bioactive coating can be an attractive solution to reduce the release of

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metal ions from stainless steels. Surface coatings change the biomaterials' surface composition, structure and morphology, retaining their bulk mechanical properties as such [5]. The concept of applying HA on metallic implants as a bioactive coating was developed due to its chemical composition close to that of human bone. The theoretical Ca/P ratio [(Caio(P04)6(OH)2), Ca/P=5/3] should be approximately 1.67 according to the stoichiometry of standard HA [6]. HA stimulates stronger bonding between the human tissue and metallic implant. Commonly used HA coating techniques are sol-gel [7], ion implantation [8], chemical vapor deposition [9], laser deposition [10], sputtering [11] and thermal spraying [12]. In this study, a thermal spraying technique was used to deposit HA coatings on 316L SS in order to ascertain its role on me corrosion behavior of the steel. The HA coatings were characterized by XRD, SEM and EDS techniques. The corrosion behavior has been investigated by Tafel extrapolation and cyclic polarization methods in simulated body fluid (Ringer's solution).

Materials and Methods Feedstock Powder and Substrate The average particle size of HA powder used in this study was 30 urn [Captai 30, Plasma Biotal Ltd. (UK)]. Commercially available biomédical stainless steel (316L SS) with chemical composition (in weight %)-C: 0.024; Cr: 16.850; Ni: 10.735; Mo: 2.269; Si: 0.468; Mn: 1.156; P: 0.032; S: 0.017 and Fe: balance was used as the substrate in this study. 316L SS coupons, each measuring 20mm x 15mm x 2mm were prepared. The specimens were polished by different grades of silicon carbide papers followed by cloth wheel polishing with alumina paste, on a polishing machine. These specimens were cleaned in acetone followed by air drying. The substrates were then grit blasted using alumina (AI2O3) grits.

Development of Coatings The 316L SS specimens were thermal sprayed using flame spray system (CERAJET) at Metalizing Equipment Company Private Limited (MECPL), Jodhpur, India. In this system, acetylene and oxygen were used as combustion gases and air was used as the carrier gas to feed the HA powder from hooper to the spray torch. This system is a high velocity version of flame spraying system and is specially designed for ceramic coatings. The spraying parameters used for HA coatings are given in Table 1.

Table l.Thermal Spray Process Parameters for HA Coatings Spraying Parameter Acetylene flow rate( 1 min"1) Oxygen flow rate( 1 min'1) Air pressure( kg cm"2 ) Powder feed rate (g min"') Spray distance (cm)

Value 73 44 4.5 15 10

Characterization of Powders and Sprayed Coatings Coated samples were characterized using XRD and SEM/EDS techniques. The phase composition of the starting powder as well as the as-sprayed coatings was analyzed by an x-ray diffractometer (X'pert-PRO, Make Netherland) with Cu-Ka radiation, operating at 40 KV/30 mA. The coated samples were scanned over 26 range of 20-60*. The morphology of feedstock powders and HA coatings was examined by SEM (QUANTA 200 FEG, Make: FE1 Netherlands) coupled with EDS. After surface characterization, the samples were sectioned with a low speed precision saw at 75 rpm speed and mounted in epoxy resin using specimen mounting press. Mounted samples were polished with emery papers of 220, 400, 600 and 800 grade and subsequently 1/0, 2/0, 3/0 and 4/0 grades. Then samples were grounded and mirror polished with

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slurry of alumina on a napped cloth to highlight the surface of HA coated substrates. The samples were gold plated to achieve the desired conductivity before observation in scanning electron microscope. Surface as well as cross-sectional SEM/EDS analysis of the coated samples was done to study their elemental compositions along with micro structural features.

Electrochemical corrosion studies To investigate the electrochemical corrosion behavior of the uncoated and HA coated 316L SS specimens, potentiodynamic polarization tests were conducted using a Potentiostat/ Galvanostat (Series G-750; Gamry Instruments, Inc. USA), interfaced with a computer and loaded with Gamry electrochemical software DC 105. The electrolyte used for simulating human body fluid conditions was Ringer's solution with chemical composition (in g/1) as 9 NaCl, 0.24 CaC12, 0.43 KC1 and 0.2 NaHC03 at pH 7.2. Before conducting the corrosion studies, each specimen was immersed in Ringer's solution for 24 hours for stabilization. The temperature of Ringer's solution was kept at 37 ± 1°C, as this is the normal temperature of human body. The temperature was maintained using a heating mantle. The exposed area of the samples in the Ringer's solution was 1cm2. The 316L SS specimen forms the working electrode. All potentials were measured with respect to a saturated calomel electrode (SCE) as the reference electrode. A graphite rod served as the counter electrode. All the tests were performed at a scan rate of lmV/sec and fresh solution was used for each experiment. Polarization curves were initiated at -250 mV below the open circuit potential and the tests were started after a steady open circuit potential was achieved (not more than ± 5mV drift in 5 min). The corrosion rate was determined using the Tafel extrapolation method and all the tests were carried out on three fresh samples to verify reliability of test results. Cyclic polarization tests were also performed to compare the pitting corrosion resistance of the HA coated and uncoated 316L SS specimens in Ringer's solution at 37 ± 1°C.

Biocompatibilitv studies In order to evaluate the biocompatibility of the HA coatings, cell culture studies were carried out. Osteosarcoma cell line KHOS-NP (R-970-5) supplied by NCCS (National Centre for Cell Science, Pune, India). These cells were seeded on three surfaces, polystyrene culture plate (control), HA coated 316L SS (H A/SS) and uncoated 316L SS (SS) in culture medium in 12 well plates at a density of lxlO3 cells/well. The culture medium consisted of 88 ml Dulbecco's modified essential media, 10 ml fetal bovine serum and 1 ml non essential amino acids. These samples were incubated for 7 days at 37°C in a humidified incubator in the presence of 5% CO2. Media was changed after every three days. After 3 and 7 days of incubation the cell proliferation was studied by methyl thiazole tetrazodium (MTT) assay. The absorbance was measured using UV-VIS spectrophotometer (Specord 250 plus, Analytik Jena, Germany).

Results and Discussion Phase analysis using X-rav diffraction analysis The XRD scans of the HA powder and the thermal spray HA coating are shown in Fig. 1(a) and (b) respectively. The scan indicates that the major peaks belong to HA (JCPDS card no. 9-432). The analysis for the HA coating confirmed the presence of HA, with minor peaks for TTCP (Ca4(P04)20) and TCP (Ca3(P04)2) phases.

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Figure 1: X-ray diffraction pattern of (a) HA powders and (b) Thermal sprayed HA coating on 316L SS, where ß-TCP (ß), TTCP (T) and HA (unmarked peaks).

SEM/EDS Analysis The SEM analysis of the feedstock HA powder confirmed that all the powder particles were spherical and had nearly uniform size. SEM micrograph of HA coating (Fig. 2) shows partially melted and un-melted particles including spherodized particles in well-flattened splats like microstructure. The microstructure consists of fiilly molten splats which are fused to each other to give a dense appearance. There is a presence of some small globular particles in the matrix which most probably are the un-melted particles. The EDS point analysis (Fig. 2) of HA coatings confirms the presence of Calcium (Ca), Phosphorous (P) and Oxygen (O) elements, which are main elements of HA. This elemental identification with their atomic % value provides information about the Ca/P ratio in the HA coating. The Ca/P ratio confirms the presence of different phases in HA coating. The Ca/P ratio of 1.64 (at point 2) is characteristic for HA which is very close to the theoretical Ca/P ratio of 1.67. The Ca/P ratio of 1.52 and 1.54 (at points 3 and 4) indicates the formation of tri-calcium phosphate (TCP). The presence of both HA and TCP phases is also confirmed by the XRD analysis.

Figure 2: FE-SEM/EDS analysis composition of thermal-sprayed HA coating on 316L SS.

Cross-sectional SEM/EDS analysis of the HA coating on the 316L SS is shown in Fig. 3. The average value of the coating thickness measured from cross-section micrograph is found to be 140 p.m. The coating has a typical splat like laminar morphology. The interface between the

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substrate and coating seems to be defect-free. The EDS peaks indicate that the HA coating contains Ca, P and O elements. The peaks of Fe, Cr, and Ni are appurtenance to 316L SS. EDS spectrum confirmed that no peak corresponding to any impurity element is present in both HA coating and 316L SS substrate.

Figure 3: Cross-sectional SEM/EDS analysis of thermal sprayed HA coating on 316L SS.

Corrosion Behavior The potentiodynamic curves of un-coated and HA coated SS in Ringer's solution at 37±1C temperature are shown in Fig. 4. The corrosion parameters determined from these curves by Tafel extrapolation method are summarized in Table 2. The various corrosion parameters are anodic tafel slope (ßa), cathodic tafel slope (ßc), corrosion potential (Ecorr) and corrosion current density (Icon-)- The corrosion potential (Ecorr) and corrosion current density (Icorr) are given by inter-section point of the cathodic tafel and the anodic tafel slope. It is well known that corrosion current density (ICOIT) is representative of the degree of degradation of metallic material. The higher the corrosion current density (lew) at a given potential, the more prone is the material to corrode. It was reported that the magnitude of ECorr is not a parameter that allows characterization of the corrosion phenomenon in a given system. Its magnitude depends on various factors such as temperature, pH and surface state of the metal [12]. According to the analysis of Tafel slope values, the results show that corrosion current density

of uncoated 316L SS in Ringer's solution (Icorr= 1.010 uAcm"2) is higher than the HA coated samples (Icorr= 193.0 nAcm"2). It indicates the susceptibility of the bare 316L SS towards corrosion attack after immersion in Ringer's solution. Moreover, the low Icorr values for the HA coated samples indicate that the HA coating prevent the substrate surface from corrosion attack during their immersion in Ringer's solution. It means that the HA coated samples (Icorr= 193.0 nAcm"2, Ecoir= -485 mV) are more corrosion resistant than uncoated 316L SS (Icorr= 1.010 uAcm2, Econ= -127 mV) in Ringer's solution. Anodic tafel slope (ßa) and cathodic tafel slope (ßc) values also shows the tendency of uncoated 316L SS towards higher corrosion. Cyclic polarization curves (Fig. 5) of HA coated and uncoated 316L SS indicated that pitting corrosion resistance of the coated sample is higher. The size of pitting loop is indication of pitting corrosion and the larger loop area shows the greater tendency to pit. These results are consistent with previous corrosion studies on HA coatings [13-14], The two layer coating composed of HA/Ti on SS316L was reported to be more corrosion resistant than the single HA layer on SS316L [14]. In another In-vitro corrosion study [15], it was reported that post heat-treated plasma-sprayed HA coatings showed much higher corrosion resistance than as-sprayed HA coatings.

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Figure 4: Potentiodynamic curves of (1) Uncoated and (2) Thermal spray HA coated 316L SS in Ringer's solution at 37±1°C.

Figure 5: Cyclic polarization curves of (1) Uncoated and (2) Thermal spray HA coated 316L SS in Ringer's solution at 37±1 °C.

Table 2. Corrosion Parameters of Uncoated and HA Coated 316L SS Parameters ßa(e'3V/decade) ßc(e"3V/decade) Ecorr(mV) ICorr(uAciïf2)

Uncoated 205.4 257.3 -127.0 1.010

HA coated 256.4 221.4

- 485.0 0.193

Biocompatibility studies

In this work, cytotoxicity tests of HA coated 316L SS (HA/SS) and uncoated 316L SS specimens were performed using osteosarcoma cell lines. The result of MTT assays are shown in Fig. 6. The absorbance values are proportional to the number of viable cells. MTT assay showed that the adherent cells proliferated on all the investigated specimens. HA coated 316L SS (HA/SS) induced a significant increase in cell viability than uncoated 316L SS samples after 3 and 7 days of culture. An implant with higher biocompatibility gives an increased cell viability which results in better biological fixation of implant [16].

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Figure 6. MTT assay results after (a) 3 days and (b) 7 days of incubation in culture medium.

Conclusions In the present study, the corrosion behavior of uncoated and HA coated 316L SS implant material was evaluated by electrochemical techniques. The HA coating was found to be useful to enhance the corrosion resistance of the steel. The following conclusions have been drawn from the study:

XRD analysis revealed that HA coatings were consisting of crystalline HA with minor TCP and TTCP phases. A coating thickness of 140 um could be achieved. The splat-like HA coating was found to have defect-free interface with the substrate steel. EDS spectrum confirmed that no peak corresponding of any impurity element was present in both HA coating and 316L SS substrate.

Based on the experimental results of Tafel and cyclic polarization, HA coated 316L SS exhibited better corrosion resistance than uncoated 316L SS in Ringer's solution. It can be deduced that HA coating prevents the metal ion release from 316L SS into human body and also decrease its susceptibility to pitting corrosion, providing a great potential in biomédical applications. It shows that HA coatings could be a viable alternative for improving the corrosion resistance of 316L SS.

The cell culture studies indicated that HA coated 316L SS demonstrated higher cell viability than uncoated 316L SS after 3 and 7 days of incubation. However, further in-vivo studies of these thermal sprayed HA coatings are necessary in order to confirm these in-vitro studies.

Acknowledgments The work presented in this paper has been supported by ISIRD Research Grant from Indian Institute of Technology Ropar, Roopnagar, India. The authors are grateful to Metalizing Equipment Company Private Limited (MECPL), Jodhpur, India for providing Flame spraying (CERAJET) facility to deposit the coatings.

References

[1] S. Yilmaz et al., "The effect of bond coat on mechanical properties of plasma-sprayed A1203 and A1203 -13wt% Ti02 coatings on AISI 316L stainless steel," Vacuum, 77(3) (2005)315-321.

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[2] A. Balamurugan, S. Karman, and S. Rajeswari, "Evaluation of Ti02 coatings obtained using the sol-gel technique on surgical grade type 316L stainless steel in simulated body fluid," Materials Letters, 59 (24-25) (2005), 3138-3143.

[3] A. Parsapour et al., "The effect of surface treatment on corrosion behavior of surgical 316L SS stainless steel implant, International Journal of ISSI," 4(1-2) (2007), 34-38.

[4] M. Sivakumar and S. Rajeshwari, "Investigation of failures in stainless steel orthopaedic implant devices: pit-induced stress corrosion cracking," Journal of Materials Science letters, 11 (1992), 1039-1042.

[5] T. Hanawa, "Roles of Metals on Regenerative Medicines," Journal of Hard Tissue Biology, 14(2) (2005), 140-142.

[6] Q. Yuan and T. D. Golden, "Electrochemical study of hydroxyapatite coatings on stainless steel substrates," Thin Solid Films, SIS (2009), 55-60.

[7] A. Balamurugan, S. Kanan, and S. Rajeswari, "Bioactive Sol-Gel Hydroxyapatite Surface for Biomédical Applications - In Vitro Study," Trends in Biomaterials and Artificial Organs, 16(1) (2002) 18-20.

[8] D. Krupa et al., "Effect of Dual Ion Implantation of Calcium and Phosphorus on the Properties of Titanium," Biomaterials, 26(16) (2005), 2847-2856.

[9] J.A. Darr et al., "Metal Organic Chemical vapour Deposition (MOCVD) of Bone Mineral like Carbonated Hydroxyapatite Coatings," Chemical Communications, 6 (2004), 696-697.

[10] V. Nelea et al., "Pulsed Laser Deposition of Hydroxyapatite Thin Films on Ti-5A1-2.5Fe Substrates with and without Buffer Layers," Applied Surface Science, 168 (1-4) (2000), 127-131.

[11] J.Z. Shi, et al., "Application of Magnetron Sputtering for Producing Bioactive Ceramic Coatings on Implant Materials," Bulletin of Materials Science, 31(6) (2008), 877-884.

[12] S. Sobieszczyk, "Surface Modifications of Ti and its Alloys," Advances in Materials Science, 10(1) (2010), 29-42.

[13] S.R. Sousa and M.A. Barbosa, "The effect of hydroxyapatite thickness on metal ion release from stainless steel substrates," Journal of Material Science: Materials in Medicine, 6 (1995), 818-823.

[14] M.H. Fathi et al., "In Vitro corrosion behavior of bioceramic, metallic and bioceramic-metallic, coated stainless steel dental implants" Dental Materials, 19 (3) (2003), 188-198.

[15] Y.P. Lee et al., "In vitro characterization of postheat-treated plasma-sprayed hydroxyapatite coatings," Surface & Coatings Technology, 197 (2005), 367-374.

[16] K.K. Saju et al., "polycrystalline coating of hydroxyapatite on TiAl6V4 implant material grown at lower substrate temperatures by hydrothermal annealing after pulsed laser deposition," Engineering in Medicine, 223 (2008), 1049-1056.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MODIFICATION RESEARCH ON THE INFLUENCE ON CORROSION

FILM PROPERTIES OF Pb-Ca-Sn ALLOYS OFWITH ADDITION OF Bi,

Ag and Zn

Lei Xu1, Lijun Liu1, Peixian Zhu2

'Faculty of Metallurgical & Energy Engineering, Kunming University of Science and Technology, Kunming, 650093, PR China 2Facuity of Materials Science and Engineering, Kunming University of Science and Technology, Kunming 650093, PR China

Keywords: Pb-Ca-Sn Alloy; cyclic voltammetry; corrosion film; electrochemical property

Abstract

The influences of addition of Bi, Ag and Zn on the electrochemical corrosion film properties of Pb-Ca-Sn(Ca=0.08%, Ca=1.2%) alloys were studied in this paper. The electrochemical properties of corrosion film were investigated by XRD, SEM, Tafel curve test and cyclic voltammetry (CV) test, respectively when the Pb-Ca-Sn alloys were in 1.28 g/mL of H2SO4 solution. The results show that effects of different elements on the electrochemical corrosion film properties of Pb-Ca-Sn alloys are different and obvious. More specific, the corrosion film becomes thicker and the crystal particles are bigger when Bi is added, and the corrosion product is mainly PbO; However, the corrosion film becomes thinner and crystal particles are smaller when Ag and Zn are added, and the corrosion product is mainly PbC>2.

Introduction

The research on Pb-Ca alloy dates back to 1859. In 1935, Haring and Thomas conducted a series of experiments and proved that lead calcium alloy showed good mechanical properties and high hydrogen evolution over-potential, thus it could be used for the production of lead-acid battery, since then Pb-Ca alloy was officially used in lead-acid batteries. However, due to its poor performance in deep discharge ability and short life, it was only used in floating charge initially. In the following 40 years, its alloy compositions, especially the content of trace element didn't show an obvious change. It was not until the 70s of last century when demand for less maintenance and maintenance free batteries increased gradually, that Pb-Ca alloy was widely studied and used. But Pb-Ca alloy still has its shortcomings [2,3], for example, it was easily to form high impedance anodic film in the alloy surface, which could lead to a dramatic decrease in the film's conductive performance, and the binding capacity between grid and active material

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also decreased, consequently, the battery's deep charge and discharge cycle performance was affected seriously. Although the growth of high impedance corrosion layer can be restricted by adding proper amount of Sn, but the effect is not ideal [4,5]. Therefore, searching for new alloy additives that can reduce the impedance of lead anodic films effectively is still the hotspot in the field [6,7].

In the present study, in order to improve the conductivity of high impedance oxide layer formed in the Pb-Ca alloy and enhance the mechanical strength of the alloy, Pb-Ca-Sn alloy was used as the mother alloy, and Bi, Ag and Zn were added to the alloy respectively, and then the electrochemical properties of corrosion film were investigated by XRD, SEM, Tafel curve test and cyclic voltammetry (CV) test, respectively when the Pb-Ca-Sn alloys were in 1.28 g/mL of H2SO4 solution.

Experimental Materials

Three different kinds of cast alloy were used in the study, the compositions of alloies were as follows:

Alloy A: Pb-0.08%Ca-1.2%Sn-0.06%Bi, Alloy B: Pb-0.08%Ca-1.2%Sn-0.03%Ag, Alloy C: Pb-0.08%Ca-1.2%Sn-0.06%Zn.

XRD and SEM Sampling and sampling set were conducted to each alloy respectively, then polish each

sample surface with metallographic sandpaper, then conduct polishing treatment to each alloy, until the surface is smooth and scratchless. And then, put the electrode disposed in series in the 1.28 g/ml H2SO4 solution for anode polarization corrosion treatment. The polarization current density is 0.5 A/cm2, polarization time 3h. After the poling treatment samples were dried, they were observed by scanning electron microscopy. After the alloy electrode corrosion film appearance were observed by the SEM, X ray diffraction analysis was conducted.

Tafel Curves and Cyclic Voltammetry Test The alloy were cast according to the composition, then roll them into same-thickness (1mm)

sheets, and process the sheets into samples of the same size and shape, the specifications of which is 2 X 1 cm, copper wire came from one side. The surface and profile of the sample was sealed with wax, an 1 cm2 area was left for reaction. Tafel curve test instrument utilize the electrochemical testing system developed by Huazhong University of science and technology. The sample surface was polished smooth with metallographic sandpaper before test and the sample received oil removal treatment. As in the electrolyte of 1.28 g/ml H2S04, a scanning speed is 0.02 mA/s, the reference electrode is calomel electrode. Auxiliary electrode is pure lead. The cyclic voltammetry utilize CHI660B electrochemical workstation and the scanning speed is 50 mV/s.

Results and Discussion

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SF.M Analysis of Corrosion Film The corrosion samples after magnified 5000 times were tested by SEM and the results were

shown in Fig.l. The results show that the morphology and structure are quite different after the corrosion film was added different elements. Specifically, Alloy A (with 0.06% Bi) has large-crystal-grain and thicker corrosion films. The grains have a close combination between each other. While alloy B (with 0.03% Ag) and alloy C (with 0.06% Zn) have small and uniform crystalline grains, which indicate that the Pb-Ca-Sn alloy has thinner corrosion films when Ag and Zn are added.

Fig.l SEM images of different alloys corrosion film

XRD analysis of Corrosion Film The conductivity of corrosion films are mainly determined by their composition. So it is

important for understanding the conductivity of corrosion films to determine the corrosion film composition. X ray diffraction analysis technique is an important method to determine material composition. In this paper, X ray diffraction analysis technique was used to study the composition change after different elements were added into the Pb-Ca-Sn alloy. The alloy XRD test results are shown in Fig.2.

The XRD results show that Alloy A (with 0.06% Bi) contains large amounts of PbO and PbOn ( n < 1 ) in its corrosion film. It is a fact that the corrosion film shows poor conductivity when the corrosion film generated non-stoichiometric oxides PbOn and n<l. As can be seen, when Bi is added into the Pb-Ca-Sn alloy, the corrosion film may generate large amounts of lead oxide with poor conductivity. While alloy B (with 0.03%Ag) and alloy C (with 0.06%Zn)

89

contain large amounts of Pb02, which has good conductivity. Therefore, the addition of Ag and Zn elements can improve the conductivity of alloy corrosion film. Some researchers also think, the addition of Zn element is not only beneficial to the forming of Pb02 film, but also restrict the forming of Pb[II] oxide film, thus enhancing conductance of Pb[II] oxidation corrosion layer and reducing its impedance.

Fig.2 XRD of different alloys corrosion film

Tafel Curve Test Fig.3 shows that Tafel curve test results of each alloy in 1.28g/mL H2SO4 solution.

Fig.3 Tafel curve of different alloys

90

The linear zone corrosion current density calculation is based on Stern-Geary equation:

^ ^ x ± - (1)

While the processing of weak polarization region data .which is based on the principles of corrosion electrochemistry, used Gauss-Newtonian and wheat Quart iterative method for curve fitting. The polarization curve equation of Weak polarization can be expressed as follows:

exp —- —

Where R is polarization resistance, i is tracking polarization current density, i is

corrosion current density, i is the limiting diffusion current density, AE = E - E0 is

polarization potential, ß , ßc are anode and cathode Tafel slope respectively. Because there

are four parameters ( i , i , ß and ß ) in formula (2), it is also known as the four parameters polarization curve equation [10,11]. In this paper, CorrTest electrochemical measurement system is used for conducting linear fitting and four-parameter fitting of each alloy's Tafel test results. The fitting results are shown in table 1.

Alloy

A B

C

Table 1 Results of fitting about Tafel curves R P " c m " 2 9m V

782.2 -0.359 1013.0 -0.451 1018.0 -0.345

with different alloys

icm mAcm"2

0.0649

0.0480 0.0312

As can be seen from table 1, the corrosion current density of alloy A which is added Bi, is stronger than that of alloy B and C. This indicates that after 0.06%Bi is added into the Pb-Ca-Sn alloy, the corrosion resistance of the alloy is weaker than alloy B and C which are added 0.03%Ag and 0.06%Zn respectively. According to formula (1), : and R show an inverse

con R p

ratio, this means that the larger the corrosion current : is , the smaller Polarization resistance con

of n will be. With the increase of j value, the corrosion process aggravates and metal

dissolves faster, which leads to the decrease of ™ polarization resistance.

The Cyclic Voltammetrv The cyclic voltammetry curves of each alloy are shown in Fig.4. The test results show that,

91

for lead alloy of different composition, their anodic process shows obvious oxidation and reduction peak, but there are no other special reaction peaks. Therefore, adding different metal elements do not change the electrochemical behavior of the grid alloy. These compositions haven't significantly changed the oxidation and reduction potentials of PbC>2. The alloy cyclic voltammetry results are shown in table 2.

As can be seen from table 2, alloy C and A have a high reduction peak, while alloy B has the smallest reduction peak value, which indicates that alloy A and C have stable surface oxide and won't be reduced easily. In addition, as can be seen from the oxidation peak, alloy A and B have higher oxidation peak, while alloy C has the smallest oxidation peak, which indicates that alloy C can be easily oxidized, and alloy A and B won't be easily oxidized.

Fig.4 Cyclic voltammetry curve of different alloys

Table 2 Results of cyclic voltammetry experiments Alloy

Oxidation peak potential EPN

Oxidation peak current Ip/A

Reduction peaks potential EPN Reduction peak current Ip/A

A

-0.82 8.14xl0"3

-0.705 2.18xl0"2

B

-0.84 7.90x10"3

-0.725 1.42xl0"2

C

-0.775 4.80xl0-3

-0.68 2.52xl0-2

Combine it with the former X ray diffraction test results analysis, after adding 0.06% Zn element into Pb-Ca-Sn alloy, the alloy has higher reduction peak current, the main product Pb02 in the corrosion film can exist stably, thus improving the alloy corrosion film's conductive properties, and at the same time it plays a positive role in prolonging lead-acid battery's life cycle. This has the same view with Literature [12]. However, alloy B with 0.03%Ag has a lower reduction peak current value, therefore a stable PbÛ2 can't be formed. But redox reversible reaction can be carried out easily in alloy electrode, which enables the electrode to have better recycling properties.

In addition, as can be seen from it, after adding 0.06%Bi elements into Pb-Ca-Sn alloy, the corrosion film generate large amounts of PbO and PbOn ( n < 1 ), which are of poor conductivity.

92

They have high reduction peak and can exist stably in the corrosion film. And they will accumulate when the electrode are used continuously, which increase the impedance of the corrosion film and thus have a detrimental influence on the cycling use of the alloy electrode. About the influence that Bi has on lead calcium electrode's performance, there is a big controversy. Some researchers [13,14] believe that Bi has a good effect on the electrochemical properties of lead calcium alloy, and it can also improve the corrosion resistance of lead calcium, reduce the release of gas when the battery is charging prolong the cycle life of the battery. But there are also some studies show that [15, 16], the addition of Bi will reduce the alloy's overpotential for oxygen evolution and aggravate corrosion or lead to the active material abscission. After several experiments, this study suggests that, adding Bi into lead calcium alloy will have a detrimental influence on the battery's cycling performance.

Conclusion

The present study investigated on the corrosion film structure, composition and electrochemical performance of Pb-0.08%Ca-1.2%Sn alloy after adding Bi, Ag, and Zn elements. The following conclusions are drawn as follows:

Scanning electron microscope results showed that the alloy corrosion film with the addition of 0.03%Ag and 0.06%Zn had fine crystallization particles and thin film. While the alloy corrosion film with the addition of 0.06%Bi had large crystallization particles.

Tafel curve test results showed that the alloy with the addition of Ag and Zn had better corrosion resistance.

XRD and cyclic voltammetry test results showed that the alloy with 0.06% Ag and Zn, with high content of Pb02 in the corrosion film had good cycling performance. But alloy with the addition of 0.06%Bi had poor cycling performance due to that it contained large amounts of PbO and PbOn ( n < 1 ), which had a poor conductivity, thus increasing the impedance of the corrosion film.

References [1] R.J.Ball,R.Evans,M.Deven, et al., "Characterisation of Defects observed within the Positive

Grid Corrosion Layer of the Valve Regulated Lead/Acid Battery", Power Sources, 2002, no.l03:207-212.

[2] H.Liu, Q.Zhou Wang, W.Zhou, "Reconsideration of some Fundamental Aspects of Anodic Pb (H) Films on Lead and Its Alloys in Sulfuric Acid Solution", Power Sources, 1999, no.84:107-113.

[3] J.L.Weininger, E.G.Sinek, "Corrosion of Lead Alloy at High Anodic Potentials", Electrochem.Soc, 23(5) (1976), 602-606.

[4] Hou-Tian Liu, Chun-Xiao Yang, Hai-He Liang, "The Mechanisms for the Growth of the Anodic Pb(II) Oxides Films formed on Pb-Sb and Pb-Sn Alloys in Sulfuric Acid Solution" Power Sources, 2002, no. 103:173-179.

93

[5] P.Simon, N.Bui, N.Pebere, et al., "Characterization by Electrochemical Impedance Spectroscopy of Passive Layers formed on Lead-Tin Alloys in Tetraborate and Sulfuric Acid Solutions", Power Sources, 1995, no.55:63-71.

[6] C.X. Yang, et al., "Impedance Study of Anodic Pb ( II ) Film Formed on Pb-Gd and Pb Electrodes in Sulfuric Acid Solution", Journal ofFudan University (Natural Science), 43(4) (2004), 584-588.

[7] J. Yang, et al., "Comparision of the Anodic Films Formed on Pb-Ce Alloy and on Common Grid Alloys in Sulfuric Acid Solution", Journal ofFudan University (Natural Science), 39(4) (2000), 427-431.

[8] Q. R. Zhu, Lead battery technology (Beijing, China Machine Press, 2002). [9] F.H. Wu, et al., " Effect of znic content on the electrochemical behavior of Pb-Zn alloys", Storage Battery, 42(2) (2006), 51-55. [10] C.N. Cao, Principles of Electrochemistry of Corrosion (Beijing, Chemical Industry Press,

2004). [11] Y.H. Liu, Electrochemical Measurement Technology (Beijing, Beijing Aviation Institute

Press, 1987). [12] W.B. Lin, R.D. Li, "The application Of Pb-Ca-Sb-Zn alloy to VRLA battelries", Storage

Battery, 1996(4): 18-20. [13] L.T.Lam, N.P.Haigh, D.A.J.Rand, "Understanding the Mechanism by which Bismuth

Improves Lead-Acid Battery Capacity", Power Source, 2000, no.88:11-17 [14] N.E.Bagshaw, "Effect of Cobalt in Lead/Acid Batteries", Power Sources, 1997, no.64:

91-95. [15] L.T.Lam, H.Ceylan, N.P.Haigh, et al., "Influence of Bismuth on the Charging Ability of

Negative Plates in Lead-Acid Batteries", Power Source, 2002, no.107: 155-161. [16] S.Zhong, J.W ang, H.K.Liu, et al., "Influence of Bismuth on Hydrogen and Oxygen

Evolution on Lead-Calcium-Tin-Aluminium Grid Alloys", Power Sources, 1997, no.66:159-164.

94

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EVALUATION OF RESIDUAL STRESS IN Fe2B COATING ON DUCTILE CAST IRON

M. A. Dofiu-Ruiz ', N. Lopez Perrusquia', V. Jorge Cortez Suarez2, D. Sanchez Huitron3

1 Institute Politécnico Nacional. SEPI-ESIME, Unidad Adolfo Lopez Mateos, Edificio 5, Av. Politécnico S/N, Gustavo A. Madero, C.P. 07338;

2 Universidad Autönoma Metropolitana, Av. San Pablo 180, Azcapotzalco Mexico, D.F, C.P. 02200

3Universidad Politecnica del Valle de Mexico UPVM, Grupo Ciencia e lngenieria de Materiales, Estado de Mexico.

Keywords: Residual Stress, Hard coating, FEM, Ductile cast iron

Abstract

The boriding thermochemical treatment enhances mechanical properties, depending on the condition under which it is produced the thin boride layer formed during a pack boriding process may be either Fe2B or FeB/Fe2B coatings and have influenced by the presence of residual stresses. The present study employ the finite element method for evaluate of thermal residual stress across Fe2B coatings produced on surface Ductile Cast Iron ASTM-A536 class 80-55-06, taking into account the power-pack boriding condition, the results of simulation was carried out in ANSYS environment were compared whit the experimental results were determinate by means of X-ray Diffraction (XRD). It was found that the residual stress varied regular with the heat treatment, globular graphites and the thickness of the coating, the distributions of residual stress determined in the Fe2B coatings are compressive with magnitudes ranging from -370 to -1740 MPa.

Introduction

One of the treatments that increase the service life of machine parts and tools is diffuses boron at the high temperature. The hard coatings can consist of the iron borides FeB and Fe2B consisting of columnar (acircular) crystals elongated in the surface direction were obtained on die surface secimens, depending on the boriding conditions and chemical composition (substrate) [1, 2]. The residual stress can be created due to different thermal expansions between substrate and coating during cooling and the different elastic properties, which usually cause spallation, distortion or the generation of cracks. Therefore, it is very important to evaluate the mechanism of causing residual stress and dieir distribution. The residual stresses greatly affect the properties of surface layers, such adhesion, resistance wear and resistance to corrosion. The service life of parts covered with such layers depends on the distribution and sign of these stresses. Also, in the literature report experimental method to evaluate the residual stress in Fe2B coatings [3-6]. Further, the residual stress in the constants thickness of Fe2B coating had been analyzed via method element finite (FEM) by D. Golanski et al [7]. The present work evaluate the distribution of thermal residual stresses obtained in the borided process by the X-ray Diffraction in the monophase Fe2B on surface ductile cast iron A-536 borides and was evaluate by finite element method to estimate the distribution of residual stress along of the Fe2B/susbtrate.

95

Material and methods

Boridine Process Ductile Cast Iron A-536 class 80-55-60 samples with a nominal composition of 3.60-3.90%C, 2.30-2.80 %Si, 0.10-0.30%A//i, 0.10 %/», 0.015 %S, 0.043%O. The samples were embedded in a closed, cylindrically shaped case stainless steel containing powder mixture (Ekabor I). The conventional pack bonding treatment was carried out at temperature of 1000 °C with exposure times of 4, 6 and 8 hours. Once the treatment was finished, the container was removed from the furnace and slowly cooled to room temperature. The borided samples were then cross-sectioned for metallographic preparation, and the depth of the surface layers was observed using JEOL JSM 6360 LV scanning electron microscope. It can clearly be seen from figure 1. The morphology of the boride layer is saw-toothed; the mean value of Fe2B boride layer thickness is approximately 85±15, 56±10 and 38±8um thick layers for 8, 6 and 4 h, respectively.

Figure 1. SEM image of microstructure on surface of ASTM A-536 borided ductile iron with dentricular morphology Fe2B layers formed at the temperature of 1000°C with a) 4 and b) 8

hours.

X-Rav Stress Analysis

Stress measurement, using the X-ray diffraction method, is based on the change in the interplanar spacing (strain) close the surface of the specimen material [8, 9]. The tests were carried out using XSTRESS3000 X-ray equipment with a CrK„ radiation, with a diffraction angle of 20=156.4° by means of a wavelength i=2.28 A operating at 30kV. The values of Poisson's ratio and modulus used for the test were taken from literature [10, 11]. Residual stress in square sample of ductile cast iron borided were determined as a function of two angles <p and y/, where y/ is the inclination angle of the sample surface normal with respect to the diffraction vector (were set between-45° and +45°) and <p denotes the rotation of the sample around the sample surface normal (were established in 0° and 90°) [12]. In each rotation angle, eight measurements on different y/ angles were performed. To obtain a residual stress-depth profile, the surface of the borided Ductile Cast iron was removed by electrolytic polishing, and the stresses on each 10 urn were measured. Finite element model

96

Residual stress in boride layers were evaluating by finite element method FEM, the ANSYS 11.0. package program [13] , in order to model the distribution of residual stress along the bonded layers after bonding, the used model represents a square shape with a Fe2B coatings on surface ductile Iron. The analytical model is assumed to be a perfect elastic body without plastic deformation in the whole analysis procedure. A symmetric plane parallel XY problem is chosen in order to reduce the data processing time; the meshes in the zone near the Fe2B/substrate interface are refined to improve the accuracy of calculation, see the geometry model in figure 2. The material properties of Fe2B coatings and substrate necessary for the numerical calculation of residual stress were taken from literature [10, 11], the values of the physical properties are Ec = 168 GPa, E, = 290 GPa, ac =6.5x10QC\ a^ll^ôxlO-C"1, vc=0.2 and vs=0.3 , where E is the Young modulus, a is the thermal expansion coefficient and v is the Poisson ratio. The subscripts c and s denoted the coatings and the substrate, respectively.

Figure 2. Schematic diagram of symmetric 2D solid a) model symmetric with the coating and substrate b) description of the geometry used in the finite element model and mesh in Fe2B layers

taken from figure 1.

The residual stresses were calculated upon cooling from the borided temperature 1000°C to room temperature 20°C. The simulation was performed by ANSYS finite element analysis to calculate the distributions of residual stresses in the Fe2B coating and the substrate. The four-node structural and triangular element PLANE 82 with symmetric element behavior formed an irregular mesh with total number of nodes 10098 and 5009 element. The thermal stress can be estimate at the condition plane stress by equation (1) [14]:

{"} = [»]{*"} (1) {«"} = {«} + {«*} (2)

{sa-} = \T-T^]{a"} (3)

where {<r}, [D],{«"'},{«},{«*},or", T and Trtf are the stress vector, elastic stiffness matrix , elastic strain vector, total strain vector, thermal strain vector, secant coefficient of the thermal expansion in the x direction, room temperature and borided temperature, respectively.

97

Results and Discussions Residual stresses In the figures 3 show the results of residual stress distribution obtained as a function of bonding time and rotation angle <p in three different directions of the boride growth. In the layer of continuous Fe2B, the compressive stresses decrease, and toward the boride-substrate interface, they change to tensile stresses conesponding to the steel substrate.

Figure 3. Residual stress distribution across boride coatings on ductile cast Iron at temperature of 1000°C with different time exposure.

The stresses occurring in the continuous Fe2B layer were compressive and highly anisotropic and decreased as a function of the penetration depth. Furthermore, the stresses switched from compressive to tensile toward the around the graphite and boride to ductile cast iron substrate interface. The magnitude of compressive stresses for both rotation angles observed in the near-surface region of the borided ductile iron peaked when the time exposition increased. In this case, compressive stresses observed at Fe2B layers, residual stress of compressive nature improves the service behavior of machine parts whereas such of tensile nature can be of detrimental effect. The compressive residual stress in Fe2B coatings of ductile cast iron are therefore beneficial.

Finite element model A coordinate system for the measurement, an element cube was therefore chosen within

the boride layer on a specimen of square cross section, (see figure 2 (a) ). The normal stress ax, ay

and oi served as the axial, tangential a radial stresses respectively produced in the Fe2B coatings. The figure 4 shows the contour plot of residual stress distribution near the edge of specimen after cooling room temperature. It can seen that there is remarkable stress concentration at or close the edge Fe2B coating/substrate interface.

98

Figure 4. Residual stress distribution as obtained by FEM (o> values in GPa) across the thickness of Fe2B layers, a) 4 and b) 8 h time exposure.

The figure 5 show the computer simulation result of residual stress by FEM in depth the Fe2B/substrate and are compared with the results experimental X-ray diffraction method and table I lists the results were obtained and was compared with the values obtained in the literature.

Figure 5. Distribution of the residual stress across the thickness of Fe2B coating-ductile cast iron obtained by FEM.

Table I. Comparison of the residual stresses of the Fe2B coatings on surface ductile cast iron borided.

Temperature (°Q

1000

Exposure time (h)

4 6 8

X-ray diffraction

(MPa) -974 -1345 -1754

FEM (MPa)

-980 -1355 -1740

Theoretical residual stress in borided alloy

steel (15]

-600 to -2000 MPa

99

The magnitude of residual stresses were obtained in ranging from -370 to -1740 MPa and distribution depend, to a considerable extent, on the phase composition of the boride coating, the technique of the layer production, thickness of coatings and the process parameters [16]. In the present work, the stress distribution over the layer depth depends on the selected profile and discontinues graphite and the process parameters. The formation of a jagged Fe2B coating interface due to locally high stress fields and this growth is expected to enhance the adhesion to the substrate ductile cast iron.

Conclusions

The residual stress of Fe2B coatings formed on ductile cast iron ASTM-A536 bonded, has been evaluate by finite element method package ANSYS and compared by X-ray diffraction. The computer simulation results correlate with the experimental results. The residual stresses obtained in function of the bonding time were compressive and the rotation angle tf the magnitude of residual stress increase as a function on time exposure and decreased as a function of the penetration depth. The high compressive residual stresses exist at the surface of borided ductile cast iron, which are balanced by tensile residual stresses located in graphite and depends strongly of the discontinues graphite, anisotropic and the process parameters.

References

1. S. Taktak, J. Mater. Sei., 41 (2007) 7590. 2. J.R. Davis, Surface Hardening of Steels: Understanding the Basics, ASM International,

Ohio, 1st edn 2002. 3. E. Macherauch and K. H. Kloos, Deutsche Gesellschaft fur Metallkunde-Springer, p.3,

Berlin, 1987. 4. Xi Chen, Jin Yan, Anette M. Karlsson, A 416(2006) 139-149, October 2005 5. Chatterjee-Fisher R. In; Sundarshan TS, editor. New York: Marcel Dekker; 1991. 6. R. Plummer and W: Pfeiffer, Journal of the less-Common metal, 117 (1986) 411-414 7. Golanski D. A. Marczuk, T. Wierzxhon, J. Mater. Sei., 14, (1995) 1499-501 8. G. Torten, M. Howes, T. Inoue, ASM International. 9. Totemeier, T.C. and J.K. Wrigth, Surface and Coatings Technology, 200:3955-3962,

2006. 10. A.V. Byakova, Yu. V. Milman, A.A. Vlasov, Sei. Sintering, 36 (2004) 27. 11. N. Frantzevich, F.F. Voronov, S.A. Bakuta, Handbook, First ed., Naukova Dumka Press,

Kiev, 1982 12. U. Welzel, J. Ligot, P. Lamparter, P. Vermeulen, E. Mittemeijer, J. Appl. Cryst.

38 (2005) 1-29. 13. Erdogan Madenci and Ibrahim Guven, Moaveni, 2006 by Springer Science-n, LLC 14. Luis Ortiz Berrocal, Mc Graw Hill.1998, ISBN: 84-481-2046-9 15. A. Pertek and J. Putniewicz, Mechanik 12 (1991) 441. 16. B.V. Babushkin, P.Z. Polyakov, Metallovedenie i Termicheskaya Obrabotka Metallov 7

(1973)27-30.

100

TIMIS2012 141s t Annual Meeting & Exhibition

General Poster Session

The proceedings contained in this section have not been edited or reviewed by the conference program organizers. The opinions and statements expressed in the proceedings are those of the authors only and are not necessarily those of the editors or TMS staff. No confirmations or endorsements are intended or implied.

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INFLUENCE OF HEAT TREATMENT ON THE CORROSION OF STEELS IN CCS ENVIRONMENT

A. Pfennig1, S. Schulz1, T. Werlitz1, E. Billow1 S. Wetzlich1, J. Tietböhl1, C. Frieslich1, and A. Kranzmann2

1 HTW University of Applied Sciences Berlin, Wilhelminenhofstraße 75 A, Gebäude C, 12459 Berlin

ania.pfennig@htw-berlin.de

2 BAM Federal Institute of Materials Research and Testing, Unter den Eichen 87, 12205 Berlin

Keywords: CCS, Carbon Capture and Storage, corrosion, steel

Abstract

To predict the reliability and safety during the injection of compressed emission gasses - mainly containing CO2 - into deep geological layers (CCS-technology, Carbon Capture and Storage), the influence of heat treatment on pit corrosion needs to be considered. In laboratory experiments different heat treated steels used as injection pipe with 13% Chromium and 0.46% Carbon (X46Crl3, 1.4034) and 0.2% Carbon (X20Crl3, 1.4021) as well as 16% Chromium steel X5CrNiCuNbl6-4 (1.4542) were tested. The samples were exposed up to 1 year to the distinct synthetic aquifer environment saturated with technical CO2 at a flow rate of 3 1/h. The corrosion rate generally does not exceed 0.03 mm/year. Pits with maximum pit heights around 300 urn were obtained for hardened X20Crl3 with martensitic microstructure. The least amount of pits is found on X46Crl3. The higher carbon content in, X46Crl3 (0.46% C), results in a lower amount of pits compared to X20Crl3 (0.20%).

Keywords: steel, heat treatment, pit corrosion, CCS, CO ̂ -injection, COz-storage

Introduction

Engineering a geological on-shore aquifer CCS-site (CCS Carbon Capture and Storage [1-3]) corrosion of the casing and injection pipe steels may become an issue when emission gasses, e.g. from combustion processes of power plants, are compressed into deep geological layers [4-8]. From thermal energy production it is known, that the COî-corrosion is sensitively dependent on alloy composition, contamination of alloy and media, environmental conditions like temperature, CO2 partial pressure, flow conditions and protective corrosion scales [6-16], Considering different environments, aquifer waters and pressures, the analyzed temperature regime between 40 °C to 60 °C is a critical temperature region well known for corrosion processes [17-20].

103

Figure 1. Possible formation mechanism of pits on the pipe steel X46Crl3 related to the galvanic model described by Han et al. [21] and modified by Pfennig and Kranzmann [19]

High Chromium steels exposed to CC^-environment generally precipitate slow growing pits mainly comprised of FeC03 (siderite) [4,8,17,19,21]. After the CO2 is dissolved to build a corrosive environment the solubility of FeCC>3 in water is low (PKSP = 10.54 at 25 °C). As a result of the anodic iron dissolution a siderite corrosion layer grows on the alloy surface. These reactions have been described in detail by various authors [7,17] and a precipitation model has been introduced by Han et. al [21], which was modified by Pfennig and Kranzmann [19] (figure 1): In the initial stage the steel is exposed to the corrosive environment, the CC>2-saturated brine (a), where the carbon dioxide forms carbonic acid in contact with the aquifer water (b). A ferrous hydroxide passivating film can form when its solubility limit is exceeded (c). A first reaction step

104

may be attributed to the formation of Fe[ll] compounds Fe(OH)2 [7,25], which leads to an increase of the local pH near the hydroxide film. Consequently a ferrous carbonate film may be formed (d). Then the growth of the corrosion layer will proceed internally and externally depending on the various carbon and oxygen partial pressures (e). Localized corrosion may then start when the ferrous hydroxide film is locally damaged due to mechanical or chemical effects. The highly porous non-protective ferrous carbonate is then exposed to the brine environment where the pH is lower. As a result the ferrous carbonate film dissolves and the steel is locally depassivated (f). This is accompanied by corrosion and passive film dissolution in lateral direction (g, h) followed by the detachment of the carbonate film (i). The removal of the detached film causes the pit to grow wider, because the same steps will occur from the beginning on the newly exposed surface.

This work is a follow-up of our work presented at "Air Pollution 2010" where local corrosion was proved to have a severe influence on the reliability of steels in on-shore CCS sites. The varying microstructures of differently heat treated steels used for CCVinjection into saline aquifer water reservoirs are of special interest.

Materials and Methods

Laboratory experiments were carried out to determine the dependence of heat treatment on pit corrosion behaviour. Different steels used as injection pipe with 13% Chromium and 0.46% C (X46Crl3, 1.4034) as well as 0.2% Carbon (X20Crl3, 1.4021) were tested. Also X5CrNiCuNbl6-4 (1.4542) was investigated as typical steel used for geothermal pumps. The steels were heat treated differently following protocols usual in the field of metallurgy (table 1).

Table 1 : heat treatment of samples used in exposure experiments temperature hold time cooling

material heat treatment „ _ . , , , M [°C ] [min] medium

X20Crl3 1.4021

X*SCrl3 1.4034

normalizing

hardening

hardening + tempering 1

hardening + tempering 2

hardening + tempering 3

normalizing

hardening

hardening + tempering 1

hardening + tempering 2

hardening + tempering 3

785

1000

1000/600

1000 / 670

1000/755

785

1000

1000/650

1000/670

1000/700

30

30

30

30

30

30

30

30

30

30

m

oil

oil

0Ü

OÜ

air

0Ü

0Ü

oil

oil

normalizing 850 30 air X5CrNiCuNblS-4 *

1.4542 hardening 1040 30 oil

hardening + tempering 1 1040/550 30 oil

hardening + tempering 2 1040/650 30 oil

hardening + tempering 3 1040/755 30 oil

Exposure tests in CCh-saturated aquifer brine were carried out using samples of the alloys above with 8 mm thickness and 20 mm width and 50 mm length. A hole of 3.9 mm diameter was used

105

for sample positioning. The surfaces were activated by grinding with SiC-Paper down to 120 urn under water and dipping into acetone for ca. 5 sec.. Samples of each base metal were positioned within the liquid phase [17-19], The brine (as known to be similar to the Stuttgart Aquifer [23]: Ca2+: 1760 mg/L, K2+: 430 mg/L, Mg2+: 1270 mg/L, Na2+: 90.100 mg/L, Cl": 143.300 mg/L, SO42": 3600 mg/L, HCO3": 40 mg/L) was synthesized in a strictly orderly way to avoid precipitation of salts and carbonates. Flow control (3 NL/h) of the technical CO2 into the brine was done by a capillary meter GDX600_man by QCAL Messtechnik GmbH, Munie. The pH after the experiments was between 5.2 and 5.6. The exposure of the samples between 24 h and 1500 h into corrosive CCS environment was disposed in a climate chamber according to the conditions at die geological site at Ketzin/Germany at 60 °C at ambient pressure - each material in a separated reaction vessel (figure 2). X-ray diffraction was carried out in a URD-6 (Seifert-FPM) with CoKa-radiation with an automatic slit adjustment, step 0.03 and count 5 sec. Phase analysis was performed by matching peak positions automatically with PDF-2 (2005) powder patterns. Mainly structures that were likely to precipitate from the steels were chosen of the ICSD and refined to fit the raw-data-files using POWDERCELL 2.4 [22] and AUTOQUAN ® by Seifert FPM.

Figure 2. Experimental set up of laboratory corrosion experiment.

Sample surfaces were analyzed with a light optical microscope Axiophot 2 by Zeiss. Here the kinetics was obtained by counting the pits per frame (6 frames per sample, 2 samples per parameter) and giving the average pit numer. Pit widtfis were measured light optically and pit depths were optained from 3-D-images. These 3-D-images were realized by the double optical system Microprof TTV by FRT. Mass gain was analyzed according to DIN 50 905 part 1-4. After the surface analysis the samples were embedded in a cold resin (Epoxicure, Buehler), cut and wet polished first with SiC-Paper from 180 um to 1200 um. The finishing was done with diamond paste 6 um and 1 urn. Different light optical and electron microscopy techniques were used to investigate the layer structures and morphology of the samples.

Results and Discussion

Surface and Morphology,

The complicated multi-layered carbonate/oxide structure is described in detail by Pfennig and Bäßler [17]. It reveals siderite FeC03, goethite o-FeOOH, mackinawite FeS and akaganeite Fe80g(OH)8Cli.34 as well as spinel-phases of various compositions. Carbides, such as Fe3C, were

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identified within the corrosion layer, similar to the high-temperature corrosion phenomena [24], The pits are covered with the same precipitates of the corrosion products formed on the surface elsewhere (figure 3).

Figure 3. Surface images of precipitations and pits of differently heat treated samples of 13% Cr steel X46Crl3 after 6000 h exposure time

Kinetics: corrosion rate, amount of counted pits and maximum pit depth.

To evaluate the influence of the heat treatment the samples were examined via light optical methods to predict the amount of.counted pits and the pit depths. Kinetics was obtained via weight loss according to DIN 50 905 after exposure to the CCVsaturated aquifer water. The results are given in figure 4 to 6.

Corrosion Rate: The corrosion rate generally does not exceed 0.03 mm/year and therefore is in good agreement with DIN 6601 allowing for 01 mm/year for pressure vessels. After 6000 h of exposure time all steels show an increase of corrosion rate. This may be due to the depassivation of the surface layers after long exposure to the CC>2-saturated saline aquifer environment. Since the rate is very low, the increase may as well be related to the manual descaling procedure. For samples with mora nd deeper pits a longer etching time is required possibly resulting in partial dissolution of small layers of the base material. With corrosion rates obtained via mass gain method about 0.002 mm/year X5CrNiCuNbl6-4 shows the lowest loss of base material for samples that were hardened or hardened and tempered. Normalized samples corrode around 0.01 mm/year determined after 6000 hours of exposure. The heat treatment does not influence the corrosion rate of X20Crl3 and X46Crl3. These are comparable for both steels around 0.003 mm/year. Therefore hardening or hardening+tempering X5CrNiCuNbl6-4 would provide best corrosion resistance in a CCS-site borehole in saline aquifer environment.

Local Corrosion (Pits): Pits were obtained metallographically and via optical volume measurement and are found on all 3 steel qualities with maximum pit intrusion depths around 300 urn for hardened X20Crl3 with martensitic microstructure after 6000 h of exposure. Overall figure 4 reveal that pit depths measured on X46Crl3 do not penetrate as deep as pits measured on the other steel samples. Still, the heat treatment does not influence the maximum penetration depth significantly except for hardened samples and X5CrNiCuNbl6-4 hardened+ temperingl. For the 13Cr steels (X20Crl3 and X46Crl3) normalizing and hardening+temperingl show less intrusion (8-25 urn) than the other heat treatments, while hardening+tempering 2 and 3 seem to be best for X5CrNiCuNbl6-4 (10 urn).

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Amount of Counted Pits: The heat treatment has little influence on the amount of counted pits per unit area, because there is little to no lowest amount of counted pits for one distinct heat treatment. The least amount of pits is found on X46Crl3. Comparing steels with the same chromium content of 13% the higher carbon content in, X46Crl3 (0.46% C), results in a lower amount of pits compared to X20Crl3 (0.20%). For X20Crl3 and X46Crl3 hardening and tempering2/3 samples show the lowest amount of pits after 6000 h while X5CrNiCuNbl6-4 has a rather high number of pits per m2. The lower amount of pits on different samples after 6000 h of exposure is due to surface corrosion phenomena: that is that pits consolidate to shallow pit corrosion and are not longer counted as single pits. These surface corrosion products prevent the access of corrosive media to the bulk material.

Figure 4. Corrosion rate, amount of counted pits and max. penetration depth after 6000 hours of exposure to aquifer brine water at 60 °C and ambient pressure of X20Crl3, X46Crl3 and X5CrNiCuNbl6-4 normalized and hardened prior to exposure.

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Conclusion

Pit growth cannot be calculated as easily as surface corrosion rates, because of its little predictability. Therefore it is not possible to give corrosion rates and lifetime predictions regarding pit corrosion in CCS technology. Summarizing the kinetic results the heat treatment preferred to obtain the least corrosive attack is normalizing or hardening+tempering2 for X20Crl3 and X46Crl3. For X5CrNiCuNbl6-4 the combination hardening+tempering 3 should be preferred. For heat treatments regarding hardening and tempering a significant good corrosion resistance cannot be given. Therfore long term exposure experiments and detailed microstructure analysis will be necessary.

Acknowledgement

This work was supported by the FNK (Fachkonferenz für wissenschaftlichliche Nachwuchskräfte) of the Applied University of Berlin, HTW and by IMPACT (EU-Project EFRE 20072013 2/21). The authors would like to thank B. Linke of the HTW for the helpful contribution.

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[2] M. van den Broek, R. Hoefhagels, E. Rubin, W. Turkenburg, A. Faaij, Effects of technological learning on future cost and performance of power plants with CO2 capture, Internal report: NWS-S-2008-10 (2009)

[3] GeoForschungszentrum Potsdam, CO2-SINK - drilling project, description of the project PART 1 (2006) 1-39

[4] S. Neäic, "Key issues related to modelling of internal corrosion of oil and gas pipelines - A review", Corrosion Science 49 (2007) 4308^1338

[5] S. Hurter, Impact of Mutual Solubility of H2O and CO2 on Injection Operations for Geological Storage of CO2, International Conference of the Properties of Water and Steem ICPWS, Berlin, September 8-11

[6] L. Zhang, J. Yang, J.S. Sun, M. Lu, Effect of pressure on wet H2S/CO2 corrosion of pipeline steel, No. 09565, NACE Corrosion 2008 Conference and Expo, New Orleans, Louisiana, USA, March 16th - 20th, 2008

[7] L.J. Mu, W.Z. Zhao, Investigation on Carbon Dioxide Corrosion Behaviors of 13Cr Stainless Steel in Simulated Strum Water, Corrosion Science, Manuscript No. CORSCI-D-09-00353 (2009) 1-24

[8] M. Bonis, Weight loss corrosion with H2S: From facts to leading parameters and mechanisms, Paper No. 09564, NACE Corrosion 2008 Conference and Expo, New Orleans, Louisiana, USA, March 16th - 20th, 2008

[9] J. Enerhaug, A study of localized corrosion in super martensitic stainless steel weldments, a thesis submitted to the Norwegen University of Science and Technology (NTNU), Trondheim 2002

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[II] R. Kirchheiner, P. Wölpert, Qualifizierung metallischer Hochleistungs-werkstoffe für die Energieumwandlung in geothermischen Prozessen, VDI-Berichte Nr. 2026, 2008

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[12] H. Zhang, Y.L Zhao, Z.D. Jiang, Effects of temperature on the corrosion behaviour of 13Cr martensitic stainless steel during exposure to C0 2 and Cl" environment, Material Letters 59 (2005) 3370-3374

[13] J.N. Alhajji and MR. Reda, The effect of alloying elements on the electrochemical corrosion of low residual carbon steels instagnant C02-saturated brine, Corrosion Science, Vol. 34, No. 11(1993)1899-1911

[14] Y.-S. Choi and S. Neäic, Corrosion behaviour of carbon steel in supercritical CC>2-water environments, No. 09256, NACE Corrosion 2008 Conference and Expo, New Orleans, Louisiana, USA, March 16th - 20th, 2008

[15] X. Jiang, S. Neäic, F. Huet, The Effect of Electrode Size on Electrochemical Noise Measurements of Mild Steel, 09575, NACE Corrosion 2008, New Orleans, Louisiana, USA, March 16th-20th, 2008

[16] Z. Ahmad, I.M. Allam, B.J. Abdul Aleem, Effect of environmental factors on the atmospheric corrosion of mild steel in aggressive sea coastal environment, Anti Corrosion Methods and Materials, 47 (2000) 215-225

[17] A. Pfennig, R. Bäßler, "Effect of C0 2 on the stability of steels with 1% and 13% Cr in saline water" Corrosion Science, Vol. 51, Issue 4 (2009)931-940,

[18] A. Pfennig, A. Kranzmann, "Influence of CO2 on the corrosion behaviour of AISI 420 and AISI 4140...", Air Pollution XVII, Volume 123 (2009) 409-418, ISBN: 978-1-84564-195-5, ISSN: 1746-448X

[19] A. Pfennig, A. Kranzmann, The role of pit corrosion in engineering the carbon storage site Ketzin, Germany, WIT Transactions on Ecology and the Environment, Volume 126, (2010) 109-118, ISBN: 978-1-84564-450-5

[20] http://www.standaid.nO/pronomi-3/data/f/0/01/36/9 10704 0/M-S06dlr2.pdf. "C02

corrosion rate calculation model" [21] J. Han, Y. Yang, S. Neäic, B. N. Brown, Roles of passivation and galvanic effects in

localized C02 corrosion of mild steel, Paper No. 08332, NACE Corrosion 2008, New Orleans, Louisiana, USA, March 16th -20th, 2008

[22] SW. Kraus, G. Nolze, POWDER CELL, J. Appl. Cryst. (1996), 29, 301-303 [23] A. Förster et. AI., Baseline characterization of the CO2SINK geological storage site at

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interface between steel and oxide scale in wet C0 2 containing atmospheres, NACE Corrosion Conference&Expo, Atlanta, 2009

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MICROSTRUCTURE AND PROPERTY MODIFICATIONS IN MOULD STEELS TREATED BY PULSED ELECTRON BEAMS

Kemin Zhang1

'School of Materials Engineering, Shanghai University of Engineering Science, Shanghai 201620, CHINA

Keywords: Low energy high current pulsed electron beam (LEHCPEB), Mould steels, Surface modification, Wear resistance

Abstract

Surface modifications on AISI H13 and D2 mould steels generated by the low energy high current pulsed electron beam (LEHCPEB) treatments have been investigated. From the observations of SEM, XRD and electron back scattering diffraction (EBSD) determinations, the microstructure modifications were studied in details. It is found that the formation of the metastable micorstructures in the surface layer are related to the very rapid heating, melting, solidification and cooling induced by the LEHCPEB irradiation. After the LEHCPEB treatment, the wear resistance of the mould steels can be effectively improved. This can be mainly attributed to the higher hardness of the ultra fine structures formed on the top surface and the hardened subsurface layers after treatment.

Introduction

Mould steels are widely used in industries for shaping various kinds of work pieces. Their surface properties are the important factors determining their quality and their service life time. Therefore, techniques for surface modification can be of interest to improve their performance [1-6]. By using a laser surface melting method, Chiang [7] et al. have obtained hardened layers in the subsurface of AISI H13 steel samples - which was attributed to the fine microstructures formed in the heat affected zone - but could not avoid coarse dendritic structures to develop in the top melted surface layer. Chemical modification methods such as nitriding and boriding, that unfortunately are time and energy consuming, have also been introduced to improve the properties and performances of mould steels [1, 2]. Other methods, mainly film deposition techniques [3-6], have also been used to modify the mould steels. Although the films always have higher hardness and better wear resistance they can suffer from poor adhesion with the substrate which inevitably leads to a deterioration of the properties.

The interaction of Low Energy High Current Pulsed Electron Beam (LEHCPEB) with materials has recently received enormous attention for surface treatment [8-15]. Compared with laser and ion beams [16-18] or the other previously mentioned technique, the electron beam has its own advantages [19,20]. The electron beam has a deep penetration depth in metals and it introduces pure energy into the surface layer without elemental contaminations. Its main feature is the generation of a high power density electron beam of 10 -10 W/cm with pulse duration of a few microseconds [10,14], The LEHCPEB irradiation induces dynamic temperature fields in the surface of the material, giving rise to superfast heating and melting that is followed by a rapid solidification of the very first surface layer by heat conduction towards the bulk. Concomitantly to the thermal effect, the pulse electron beam creates a dynamic stress field that causes intense deformations at the material surface [15]. The combination of these processes makes it possible to modify substantially the surface characteristics and, in many cases, improve the mechanical

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properties faster and more efficiently than conventional surface treatment techniques [21-26]. Usually, the LEHCPEB modified surface can be divided into three successive regions having different ranges of penetration depth: (i) a melted and rapidly solidified layer on the top surface (~1 um); (ii) a heat affected zone (HAZ~10 um) and (iii) a stress wave affected zone (-100 um). In the case of steels, it has been established that the hardness, wear resistance or corrosion resistance can be significantly improved by using the LEHCPEB treatment [27-29]. In particular, the microstructure present in the melted layer plays a determinant role in controlling the corrosion resistance [20-22]. Indeed, the LEHCPEB treatment and its associated rapid solidification can generate non-equilibrium solidification conditions and lead to specific features of the melted surface. These can include, for example, the formation of a supersaturated solid solution, metastable phases, ultra fine grain sizes and nano sized precipitates [8, 12,24-27].

In the present work, the effects of the LEHCPEB treatment on the microstructure and the properties of the AISI D2 and H13 mould steels have been studied.

Experimental

Typical mould steels, AISI H13 and D2, were chosen to be studied in this work. Their chemical compositions can be found in Refs. [24, 30]. Before electron beam treatments, the steel samples were machined to dimensions of 10X 10X25 mm3. The AISI H13 steel samples were heat treated at 1040°C for 30 min, water quenched and tempered at 600°C for 1 hour. The AISI D2 steel samples were heat treated at 1020°C for 30 min, water quenched and tempered at 200CC for 3 hours. Then they were ground and polished with diamond paste down to R„ « 0.07 p.m. The electron-beam treatment parameters for H13 steel sample were as follows: accelerating voltage, 27 kV; energy density, 3 J/cm2; number of pulses, 25; pulse duration, 1.5 u,s. For D2 steel samples, the parameters were set as: accelerating voltage, 18 kV (D2-1), 21.5 kV (D2-2), 27.8 kV (D2-3); number of pulses, 20; pulse duration, 1.5 us.

The phase state at the sample surface was analyzed with a Cu-K„ radiation source on a SHIMADZU XRD-6000 apparatus. A JEOL 6500F scanning electron microscope (SEM) apparatus equipped with a field emission gun (FEG) was used to observe the surface morphologies. This SEM was also equipped with an electron backscattering diffraction (EBSD) attachment which was used to gain more information about the phases and microstructures of the top surface layer.

A ball-on-disk friction and wear apparatus was used to measure the wear resistance of the untreated and LEHCPEB treated samples. The friction counterpart was a GCrl5 ball. The load was 5 N for HI3 steel samples and 3 N for D2 steel samples. For comparison, all the wear measurements of the untreated and treated samples were carried out under the same testing conditions at ambient temperature of 25 °C.

Results AISI D2 steel

Fig. la shows the typical cross-sectional optical micrograph of the treated sample D2-3 after etching. The "white" layer formed on the surface, which is corresponding to the melted layer that is known to be weakly echable [10]. The depth of the melted layer is about 4 um. No large carbides can be observed in the melted layer. In contrast, large carbides and tempered martensite lathes are present in the substrate.

To study the phase transformation during the treatment, XRD analysis was carried out on the treated and untreated samples, and the results are shown in Fig. lb. It can be seen that, the

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untreated sample mainly contains two phases, tempered martensite (a'-Fe) and carbide (M7C3). The peak intensity of the later is relatively lower, which is due to its low volume fraction. After the LEHCPEB bobardment, it is clearly that XRD patterns have changed. Firstly, some new peaks appear in the XRD patterns, they are verified as y-Fe phase and the peak intensity of the y-Fe phase increases with the beam energy (accelerating voltage), indicating that the volume fraction of the y phase increases with the beam energy. Considering that the melted depth increaes with the electron beam energy [9], the y phase should form during the resolidification of the melted layer. For sample D2-3, the fraction of y phase reaches nealy about 85% as calculated from XRD peak intensity. It is unusual to obtain such a high fraction of y phase on the surface layer, since it is a high temperature phase, and it should transform to martensite due to the very high cooling rate (~108KVs) in our case [15]. Secondly, the peaks of the carbide dissappear after LEHCPEB treatment, which is due to its dissolution as we have discussed before. Finally, all the peaks of martensite move to high angles, showing that the lattice constant reduces. As is well known, if martensite was heated to a certain temperature, when the carbon can diffuse, its lattice constant will reduce due to the carbide precipatation from martensite, and this annealing process will result in the tranformation of metastable martensite to stable ferrit. In this case, the temperture in heat affected zone is high enough for anneling the D2 steel. Therefore, a-Fe should appear in the layer between melted layer and the substrate, as we labled in the XRD patterns of the treated samples.

Figure 1. Cross section micrograph of the LEHCPEB treated D2-3 sample (a) and XRD patterns of the D2 samples (b): (1) untreated, (2) D2-1, (3) D2-2, (4) D2-3. It is well known that wear resistance is the most important property for die steels. In this case,

we studied the wear resistance of all the samples at the same experimental conditions. Friction coefficients for untreated and LEHCPEB treated D2 samples were measured and the results are shown in Fig. 2a. It shown that, for the untreated sample, the initial value of the coefficient of friction is low (about 0.25), and this may be due to the lubrication by residual contaminants. At about the 10 cycles, the friction coefficient of untreated D2 steel increased rapidly and reached the highest level 0.71, and this high coefficient of friction is probably caused by the intermolecular (adhesive) interaction between the D2 steel and counter ball surfaces. The friction coefficient variation of treated specimens is completely different. For example, the friction coefficient of D2-1 specimen maintained very low value of 0.12 for a relatively long-term from beginning to 15 cycles, and then it slowly increased to 0.7, and finally remained at this level. As to other treated samples, the same phenomena are observed. The best wear performance is found on the D2-3 sample, which remains at low friction coefficient level up to 38 cycles. The measured wear rate of the treated D2 samples reduced significantly compared with that of the initial sample. For D2-3 sample, the value is only about 1/7 of the untreated sample.

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Fig. 2b shows the wear rate versus the fraction of the y phase in the surface layer. It can be seen that, the wear rate decreases linearly with the increase of y phase content in the surface layer. The above results showed that, LEHCPEB treatment is an efficient method to improve the wear resistance of D2 steel.

Content of 7 phase

Figure 2. Friction coefficient curves of the untreated and treated samples during wear tests (a) and wear rate of D2 steel samples versus the content of y phase in the surface layer (b).

AISI H13 steel

Fig. 3a shows a typical cross section OM micrograph of the LEHCPEB terated H13 steel sample. A continuous "white" layer is visible on the surface. It has a clear interface with the substrate material and corresponds to the melted layer. The depth of the melted layer is about 2 um, which is lower than that of the D2 steel under the same treatment parameters [24, 25]. This must be a consequence of the higher heat conductivity of the H13 steel due to its lower concentrations in C and Cr. The 'V like shape observed in Fig. 3a, which has its center lower than the rest of the melted layer, witnesses the mark of an erupted crater. Below this melted layer, a needle like morphology corresponding to the initial aspect of the tempered martensite (ferrite + carbides) is still visible.

Figure 3. Cross section micrograph of the LEHCPEB treated H13 steel sample (a) and XRD patterns of the H13 samples (b): (1) untreated, (2) LEHCPEB treated.

Fig. 3b shows the XRD diffraction patterns of the untreated and LEHCPEB treated H13 steel samples. The X-ray diffraction peaks indicate that the major phase in the initial material was the a-ferrite. This phase results from the decomposition of the quenched martensite during the

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tempering process. Carbides having the MC and M23C6 structures can be also detected in the X-ray pattern. These carbides are reported to be present under the similar heat treatment conditions [31]. Because of their low volume fractions, the intensity of the carbide peaks is however rather weak. After the LEHCPEB bombardment, the observed XRD pattern has changed quite significantly. A splitting of the (200) peak is observed after 25 pulses. In addition, some new peaks corresponding to the y-Fe phase are also present in the XRD pattern. The peaks corresponding to the carbides have disappeared after the LEHCPEB treatment; a result of dissolution of carbides during treatment [25]. Compared to the results obtained in the D2 steel after 25 pulses [24], the relative intensity of the y-Fe peaks over the a-Fe ones is here much weaker. This indicates that the volume fraction of the y phase in the surface layer of this H13 steel is lower than the one obtained in the case of the D2 steel.

Figure 4. EBSD orientation map taken on the surface of the LEHCPEB treated H13 steel sample.

Figure 5. wear depth profiles of the H13 steel samples before and after the LEHCPEB treatment.

Fig. 4 is an EBSD orientation map taken on the treated H13 sample. The Kikuchi patterns detected by EBSD could be indexed by the ferrite (or martensite) and austenite phases. For the sake of clarity, the ferrite grains are shown in colour while austenite grains are shown in grey contrast in Fig. 4. The colour of the grains was selected according to the standard colour triangle given in the inset of Fig. 4. The grain boundaries are also shown in the maps. Low angle grain boundaries (LAGB) having misorientation between 2° to 15° are shown in white while high angle

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grain boundaries (HAGB) having higher misorientations are shown in black. Consistently with the XRD results, the EBSD map indicates that the ferrite (or martensite) is the dominant phase at the top surface. From the EBSD map in Fig. 4, it can be established that the range of grain size for the ferritic grains is rather fairly broad: from one hundred of nanometers up to a few micrometers.

Wear resistance of H13 samples was evaluated by measuring the depth of grinding grooves. Fig. 5 shows the curves of the wear depth as a function of time. It can be clearly seen that the slope of the curve for the LEHCPEB treated sample is much lower than that of the untreated sample. That is, the LEHCPEB irradiated sample shows an improved wear resistance compared with that of the untreated one. The fact that the wear depth profile for the treated sample does not show a linear increase with wear time may be attributed to the nonhomogenity nature of the treated surface layer and the increased surface roughness after the LEHCPEB treatment [23]. The final wear depth of the untreated sample is about 2.2 times as that of the treated sample after 20 min tests.

Disscussions

During LEHCPEB bombardment, the concentrated energy fluxes acting on the mould steel samples induced dynamic temperature fields in the surface layers to which the energy has been deposited, giving rise to superfast heating, melting, and evaporation, followed by superfast solidification in the melt zone of the material. The dynamic stress fields formed cause intense deformation processes in the material [29,32]. These processes in combination make it possible to produce in the surface layers of the materials metastable states, which may impart the material with improved physicochemical and strength properties [32-34]. The different microstructures observed on the treated D2 and H13 samples can be understood by considering the effect of the treatment on the dissolution kinetics of the carbides and the ability of the process to homogenize the melted surface. The C content in D2 steel is much higher than that in H13 steel. As a result, the austenite formed on D2 steel during the LEHCPEB treatment can be stabilized through carbide dissolution. The surface properties are determined by the final microstructure and strain-stress state of the LEHCPEB treated surface layer [10,29]. For D2 steel, the wear rate decreases linearly with the increase of y phase content. It is well known that y phase is usually softer than martensite. However, the wear resistance is not only determined by the hardness of the materials, but also the toughness. Firstly, the y phase obtained in this case directly grows from liquid during the fast crystallization process, and the dissolution of Cr and C into this phase will significantly enhance the strength of it, so it has a good combination of hardness and toughness, and high C content is also an important factor for the stabilization of the y phase [24]. Secondly, under the high compressive stress, the martensitic transformation can take place, which is observed in many other cases. So, it is reasonable that the austenite in this case will transform into martensite under the loading of the wear test, which can surely increase the surface hardness and improve its wear resistance. For H13 steel, both the top surface layer and the subsurface layers were hardened after the LEHCPEB treatment [30]. Therefore, the wear resistance improvement can be mainly attributed to the higher hardness of the ultra fine structures formed on the top surface and the hardened subsurface layer after treatment. Similar wear improvements were also observed in the carbon steels [19], and Mg alloys [23] after the LEHCPEB treatment.

Conclusions

The LEHCPEB treatment has been applied to AISI H13 and D2 mould steels and the intriguing surface modifications occurring in the melted layer have been investigated. The specific features observed in the H13 and D2 steels are compared to each other in order to improve our understanding of microstructure and property modifications associated with this technique. It is

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found that the different microstructures observed on the treated D2 and H13 samples can be understood by considering the effect of the treatment on the dissolution kinetics of the carbides and the ability of the process to homogenize the melted surface. In both cases, the wear resistance of the mould steels can be effciently improved by the LEHCPEB treatment. This is mainly due to the the ultra fine structures formed on the top surface and the hardened subsurface layers after treatment.

Acknowledgements: Prof K.M. Zhang would like to acknowledge the support of the Leading Academic Discipline Project of the Shanghai Education Commission (Grant No. J51402). This work is partially supported by the Special Foundation of the Shanghai Education Commission for Nano-Materials Research (No. 1052nm05000) and National Natural Science Foundation (No. 51101096).

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crystalline copper irradiated by intense pulsed ion beam," Surf Coat Tech. 108-109(1998), 171-180. 18. M. A. Meyers, F. Gregori, B. K. Kad, M. S. Schneider, D. H. Kalantar, B. A. Remington, G. Ravichandran,

T. Boehly, and J. S. Wark, "Laser induced shock compression of monocrystalline copper : characterization and analysis," Ada Maler. 51(2003), 1211-1220.

19. S. Z. Hao, B. Gao, A. M. Wu, J. X. Zou, Y. Qin, C. Dong, J. An and Q. F. Guan : Surface modification of steels and magnesium alloy by high current pulsed electron beam. Nucl. Instr. and Melh. in Phys. Res. B. 240(2005), 646-654.

20. J. X. Zou, K. M. Zhang, C. Dong, Y. Qin, S. Z. Hao, and T. Grosdidier, "Selective surface purification via crater eruption under pulsed electron beam irradiation," Appl. Phys. Leu. 89(2006), 041913-1-3.

21. K. M. Zhang, J. X. Zou, T. Grosdidier, C. Dong, and D. Z. Yang, "Improved pitting corrosion resistance of AISI 316L stainless steel by high current pulsed electron beam treatment," Surf Coat Tech. 201(2006), 1393-1401.

22. K. M. Zhang, D. Z. Yang, J. X. Zou, T. Grosdidier, and C. Dong, "Improved in vitro corrosion resistance of a NiTi alloy by high current pulsed electron beam" Surf Coat Tech. 201(2006), 3096-3102.

23. B. Gao, S. Z. Hao, J. X. Zou, T. Grosdidier, L. M. Jiang, J. Y. Zhou, and C. Dong, "High current pulsed electron beam treatment of AZ31 Mg alloy," J. of Vac. Sei. and Tech. A. 23 (2005), 1548-1556.

24. J. X. Zou, T. Grosdidier, K. M. Zhang, and C. Dong, "Mechanisms of nanostructures and metastable phase transformations in the surface melted layer of a HCPEB treated D2 steel," Ada Maler. 54(2006), 5409-5419.

25. J. X. Zou, T. Grosdidier, B. Bolle, K. M. Zhang, and C. Dong, "Texture and microstructure at the surface of an AISI D2 steel treated with high current pulsed electron beam," Metal. Mater. Trans. A, 38(2007), 2061-2071.

26. J. X. Zou, T. Grosdidier, K. M. Zhang, B. Gao, S. Z. Hao, and C. Dong, "Microstructures and phase transformations in the surface layer of AISI D2 steel treated with pulsed electron beam," Journal of Alloys and Compounds, 434-435(2007), 683-686.

27. Y. Qin, C. Dong, X. G. Wang, S. Z. Hao, A. M. Wu, J. X. Zou, and Y. Liu, "Temperature profile and crater formation induced by high current pulsed electron beam," J. of Vac. Sei. and Tech. A. 21 (2003), 1934-1938.

28. Y. Qin, J. X. Zou, C. Dong, X. G. Wang, A. M. Wu, Y. Liu, S. Z. Hao, and Q. F. Guan, 'Temperature-stress fields and related phenomena induced by a high current pulsed electron beam," Nucl. Instr. and Melh. in Phys. Res. B. 225(2004), 544-554.

29. J. X. Zou, K. M. Zhang, S. Z. Hao, C. Dong, and T. Grosdidier, "Mechanisms of hardening, wear and corrosion improvement of 316 L stainless steel by low energy high current pulsed electron beam surface treatment," Thin Solid Films, 519(2010), 1404-1415.

30. K. M. Zhang, J. X. Zou, T. Grosdidier, and C. Dong, "Microstructure and property modifications of an AISI H13 (4Cr5MoSiV) steel induced by pulsed electron beam treatment," Journal of Vacuum Science and Technology A, 28(2010), 1349-1355.

31. X. B. Hua, L. Lia, X. C. Wu, and M. Zhang, "Coarsening behavior of M23C6 carbides after ageing or thermal fatigue in AISI H13 steel with niobium," International Journal of Fatigue, 28(2006), 175-181.

32. T. Grosdidier, J. X. Zou, N. Stein, C. Boulanger, S. Hao, and C. Dong, "Texture modification, grain refinement and improved hardness / corrosion balance of a FeAl alloy by pulsed electron beam surface treatment under the "heating mode"," Scripta Malerialia, 58(2008), 1058-1061.

33. K.M. Zhang, J.X. Zou, T. Grosdidier, and C. Dong, "Formation and evolution of craters in carbon steels during low energy high current pulsed electron beam treatment," Journal of Vacuum Science and Technology A, 27 (2009), 1217-1226.

34. K.M. Zhang, J.X. Zou, T. Grosdidier, C. Dong, and S. Weber, "Ti surface alloying of an AISI 316 stainless steel by low energy high current pulsed electron beam treatment," Journal of Vacuum Science and Technology A, 26(2008), 1407-1414.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

POTENTIAL FIBERBOARD MATERIAL FROM COW MANURE AND DISPOSABLE WATER BOTTLE

Boon-Chai Ng1, Craig Bradfield1, Roy Pritish', and Marlene Murray2

'Engineering and Computer Science Department 2Biology Department

Andrews University, Berrien Springs, Michigan 49104, USA

Keywords: Cow Manure; disposable water bottle, fiberboard

Abstract In this pilot project, readily available solid cow manure from the nearby Andrews University dairy farm was rinsed to remove any feces, dried, and then tested for any life bacteria. This fiber material is then blended with shredded disposable plastic water bottle to form a "green' composite material. This composite material was placed in a cylindrical mold and heated to various temperatures to allow bonding of the thermoplastics to the fiber. The heated composite material was subsequently compacted with a 10,000 lbs. load using the universal tensile tester. Results showed that the composite material heated to a temperature of 250°C for an hour before compacting with a 10,000 lbs. load produced a well bonded fiberboard.

Introduction Cow manure is an excellent fertilizer for farming. It has been used for centuries as a fertilizer because it is a good source of plant nutrients and organic matter [1]. The Andrews University Dairy Farm has more than 800 cows and much of the manures are used as fertilizer in the farmland. With so much manure produced daily, some of the excess manure is sold to the local farming industries. The rest of the manure is heaped up at one corner of the land, waiting to be spread out on the farm at a later time. This causes some unpleasant odor when the wind blows from the farm towards the university. Plastic bottles are commonly used to store water for drinking outside the house. Americans, according to Container Recycling Institute (http://container-recycling.org), buy an estimated 34.6 billion single-serving (1 liter or less) plastic water bottles each year. Almost eight out often of these bottles end up in a landfill or incinerator. Hundreds of millions end up as litter on roads and beaches or in streams and other waterways. Taxpayers pay hundreds millions of dollars each year in disposal and litter cleanup costs. The large amount of cow manure from the farm and the disposable water bottles could be recycled for other uses. Cow manure can be converted to a usable fiber. In fact, Michigan State University researchers had done a pilot study of manure-based fiber. These fibers could take the place of sawdust in making fiberboard [2]. Farmers have also recycled this manure as bedding materials for their animals [3].

The objective of this project is to recycle the readily available cow manure and disposable plastic water bottle and combining them together to form a 'green' composite material. This composite material can be used as a potential fiberboard.

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Experimental Work Raw manure was collected and soaked in water for a week. The water was stirred occasionally and changed daily to remove any floating feces. The rinsed fibers were then dried in an oven at a temperature of 70°C or 120°C. Turner [4] suggested that temperature of drying the fresh manure should be above 45°C to reduce any biological activity in the manure. Samples of these dried fibers were tested for growth (live bacteria). Two different culture tests were performed on the dried fiber. The first test was inoculating the fiber with Luria Bertani (LB) broth. In this test, the test tubes were filled with the LB broth and samples of the fibers (see figure 1) and incubated overnight at 37 ° C in a rotary shaker at 200 rpm. The second test was to spread the fibers on LB agar plates (see figure 2) and incubated for 24 to 48 hours at 37° C.

Figure 1 shows some of test tubes filled with Luria Bertani (LB) broth and fibers before placing them a rotary shaker for incubation.

Figure 2 shows sample fibers placed on LB agar plates for incubation. Note the top left is the control sample.

Disposable water bottles are made from polyethylene terephthalate, PET or PETE, with a melting temperature of 250 degree C [5], These bottles were heated to 70°C, flattened and shredded with a cross cut paper shredder with a width of-3.0 mm.

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Dried fibers and shredded plastics were then blended together in a ratio of 3:2 and place in a cylindrical mold, measuring 1.0 in diameter. The cylindrical mold has inserts on both ends to contain the fiber-plastic mixture. The mold and its contents are then placed in a furnace and heated at various temperatures either 250°C or 280°C and held at that temperature for one hour or two hours before compressing the mold contents in the universal tester to 10,000 lbs.

Results and Discussion The rinsed and dried manure/fiber does not have the typical distinct smell associated with the cow manure. The culture test results showed that fibers dried in the 70°C oven had bacterial growth in both the agar plates and the broth. Figure 3 showed some of the growth that had developed around the fibers that were dried in the 70°C oven. Fibers dried in the 120°C did not show any evidence of growth. Figure 4 showed similar results for the dried fiber inoculated with Luria Bertani (LB) broth.

Figure 3 shows growth around the fibers that were dried in the 70°C oven.

Figure 4 shows test samples after incubation in the rotary shaker. The far left was fibers from the finished product and the far right is the control sample broth. The broth in these tubes is clear indicating no growth. The two middle sets of test tubes are from fibers dried in 70°C oven. Growth in these tubes makes the broth cloudy.

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Samples that were heated to 280° C for an hour had a charred appearance as showed in figure 5a. Likewise fiber/plastic mixture, soaked at 250° C for 2 hours before compaction, were also charred (see figure 5b). The most appealing condition was when the mixture was soaked in the 250° C oven for an hour before compacting with the universal tester to 10,000 lbs. The fibers did not char and the mixture was observed to bond nicely as shown in figure 5c.

Figure 5: a) Charred sample after heating at 280°C for an hour before compaction, b) charred sample after heating at 250°C for two hours before compaction, c) sample heated at 250 for one hour before compaction.

Future Work Work is still in process to continue to develop this green composite material. Some of the tests to consider include subjecting the composite material to 100% humidity conditions and examination of the de-bonding or disintegration of the fibers after the following operations:

a. Drilling onto the sample b. Pounding nails into the sample c. Putting a screw through the sample

Conclusion A green composite material had been successfully produced by blending fibers recycled from cow manure and shredded plastic bottles. This mixture was heated at 250°C for an hour before compacting with a universal tester to 10,000 lbs. The following were observed during the experiment:

• Fibers that were dried at 70°C still showed evidence of growth whereas fibers dried at 120°C did not show any growth.

• Heating the mixture to 280°C for an hour or 250°C for two hours will char the fibers. • Samples that were heated 250°C for one hour before compacting with the universal tester

to 10,000 lbs. were observed to bond nicely.

Acknowledgement The project had been funded by the Andrews University Faculty Research Grant No: 201162.

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References 1. Andy Bary, Craig Cogger, and Dan M. Sullivan, "Fertilizing with Manure", A Pacific

Northwest Extension Publication, PNW0533.

2. USA Today, 2nd July, 2007.

3. Manure Manager, July 28, 2009.

4. C.W. Turner, "Effect of Temperature and Drying of Activation of Male Hormone of Cow Manure, Journal of Dairy Science, 32:796, 1949.

5. Record of Polyethylene terephthalate in the GESTIS Substance Database from the Institute for Occupational Safety and Health of the German Social Accident Insurance (IFA).

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces IMS (The Minerals, Metals & Materials Society), 2012

INFLUENCE OF PROCESS AND THERMO-PHYSICAL PARAMETERS ON THE HEAT TRANSFER AT ELECTRON BEAM MELTING OF

Cu AND Ta

Katia Vutova1*, Veliko Donchev1, Vania Vassileva1, Georgi Mladenov'

'institute of electronics, Bulgarian Academy of Sciences, 72 Tzarigradsko shosse, 1784 Sofia, Bulgaria

*e-mail: katia@van-computers.com ; tel: +359-2-9795900; fax: +359-2-9753201

Keywords: Heat transfer, Thermal stream, Electron beam melting and refining (EBMR), Copper, Tantalum.

Abstract

Electron beam melting and refining (EBMR) in vacuum is an ecological friendly method for metal purification, scrap regeneration and reuse of expensive metals and special alloys needed for all areas of human activity. The thermal transfer processes are important for production of metal blocks with good quality (structure, surface and composition) and the mathematical modeling is a tool for studying and control of these processes. The temperature variations of the thermal conductivity and the heat capacity for Ta and Cu are estimated and are taken into account in the presented heat model. Some results on the influence of the casting velocity, beam power and ingot dimensions on the heat streams through the different boundaries and on the crystallization front shape for EBMR of Cu and Ta are presented and discussed. Calculated and experimentally obtained crystallization front forms are compared and a good correspondence is observed. Electron beam casting conditions for obtaining of good quality cupper ingot are optimized.

Introduction

Electron beam melting is a key method for obtaining new materials for micro- and nano-electronics, for which high purity and quality are demanded. This method combines well the advantages of the electron beam as unconventional source of heating without limitations on the achieved temperature and the high vacuum as an environment for the refining processes. Therefore the method is suitable for refractory metals and alloys as well as for metals and alloys, which can actively react with oxygen in air-heated environment [1-4]. Optimization and control of the complex physical, physicochemical and thermal processes during electron beam melting and refining (EBMR) are difficult and are possible if dependencies concerning the liquid metal pool and influence of the heat transfer on the crystallization processes are known. The form of the crystallization front (liquid/ solid boundary) is an important condition for producing pure metal ingots with perfect crystal structure as well as the geometry of the liquid pool influences the impurities removal. They both depend on the EBMR process parameters (beam power, casting velocity, etc.) and thermo-physical characteristics of the investigated liquid and solid metal. It is very difficult to get real-time information for the processes in the molten metal as well as for temperature distributions along the formed block [5-6]. These problems require development of models and computer tools for studying and better understanding, for optimization and successful application of EBMR for purification of different metals and

125

alloys, scrap regeneration and reuse of expensive metals and special alloys needed for energetics, medicine, transport, etc. The results presented in this work are continuation of research on the conditions and parameters influencing the thermal transfer processes in treated metal ingots at electron beam melting [7-10]. Numerical experiments are made, using the described heat model, and data for: heat flows through the different boundaries of the formed pure ingots, liquid pool geometry, crystallization front shape, etc. are obtained in the case of EBMR of Cu and Ta -metals with different melting temperatures, thermal conductivity X and heat capacity Cp. The aim is to investigate and clarify the heat transfer at different process regimes, the influence of geometrical dimentions of the ingot as well as the influence of the thermal parameters (X. and Cp), depending on the metal temperature.

Fig.l. Scheme of the electron beam "drip melting method". Gl- the top surface of the ingot; G2 - the molten ingot/ crucible wall interface; G3- the ingot/vacuum interface; G4- the interface ingot/ puller interface.

Heat Model and Thermo-physical Parameters

Heat Model

Technological scheme of the classical method - electron beam melting and refining drip process with horizontal feeder and the drip molten metal crystallized in a water-cooled copper crucible is shown in Fig.l. The temperature distribution in the metal cast ingot (R-its radius, H-its length) is described by Poisson's equation; the temperature T is a function of the radius r and the height z (Fig.2). Equations (l)-(6) define the used model [10]:

126

where: V is the casting velocity of the ingot; a is the thermal diffusivity; /V/ is the power density; Rv are vapor losses (RV=CPW, W-weight loss velocity); T, T1 and T2 are the temperatures of the two sides of the boundaries. When V^O, the model describes the solidification of the cast ingot in a crucible - "drip melting process" (Fig.l). At V=0 the model corresponds to "disks melting method" (no material is added to the ingot by pouring) [10]. The boundary conditions (Eqs. (3-6)) take into account the radiation and vapor losses and three mechanisms of heat transfer are assumed [9]. Discrete approximation of Eqs. (1-6) is made and the numerical solution is gained by simple iterations method. The stop condition is reaching a balance between the input and output thermal flows through the boundaries.

Fig.2. (a) Rectangular net of points along the ingot, (b) Approximation patterns of the equations.

Thermo-phvsical Parameters

For investigation of physical and heat processes, the knowledge of the nature of thermal parameters of the metals and alloys is an important factor. The thermal conductivity X and heat capacity Cp of metals and alloys have different values at different temperatures T and are specific for each metal and alloy. Experimental data [11,12] for the thermal parameters X and Cp is used to model these dependencies for the refractory metal Ta and low-melting-temperature metal Cu. The obtained temperature dependencies of X and Cp are presented in Table I.

Table I. Obtained functions for the thermo-physical parameters for Cu and Ta.

Metal

Cu

Thermal conductivity I, W/m.K Value

432.12-0.0937*T

320

Condition T<1300K

T>1300K

Heat capacity C0, W.s/g.K Value 0.388

0.359+0.0000996825V 0.494703

Condition T- 288K

288K <?'•- 13S6K 1356<T

Ta 75.3676880 + 0.00382799*T

82.7676880 + 0.00W2799*T

T<3000K

T-3000K

0.0000208044*T+0.1322 0.000065»T+0.0228333 0.00000644*T+0.19848

T < 2475.6 K 3000 K > T > 2475.6 K

T >3000K

Two series numerical experiments for "disks melting method" were made. The tantalum ingots are with 50 mm diameter and 20 mm length. The beam is focused in a circle with 20 mm diameter on the top ingot surface (Fig. 1). The beam power varies from 15 kW to 50 kW.

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For the experiment with constant thermal parameters the values of X and Cp are 58 [W/mK] and 0.12 [Ws/gK], respectively. The molten pool geometry and the heat fluxes are examined for both cases, when the thermal parameters are constants and when they are functions of the temperature. In Fig.3 the results concerning the depth of the liquid pools are compared. Figure 4 shows the heat flow through G4 - the bottom of the ingot (Fig. 1).

Fig.3. Molten pool depth vs. the beam power for both experiments - linear functions and

constant values for X and C„.

Fig.4. Heat losses through the bottom of the ingot vs. the e-beam power.

Calculated data shows that the difference in the results obtained at constant and non-constant values for X and Cp is not neglectable. Using functional dependencies for X and Cp> the thermal losses through the ingot bottom (G4) increase with about 2 kW for lower e-beam powers and with about 5 kW for powers higher than 30 kW (Fig.4). The obtained results concerning the liquid pool volume dependence on the beam power in the case of EBMR of Cu are shown in Fig.5. The ingot length is 100 mm and its diameter is 60 mm. The beam spot radius is 10 mm. For beam powers higher than 9 kW the difference is about 5%. The results, corresponding to the temperature dependencies for X and Cp, coincide better with the experimental results (Fig.6). Experiments on EBMR were performed using 60 kW equipment (ELIT-60).

Fig.5. Dependence of the molten pool volume on the beam power. In the series with constant heat

parameters: X=318.1 [W/mK] and Cp=0.38 [Ws/gK].

Fig.6. Liquid pool in Cu ingot. Dashed curve presents the interface liquid metal/ solid metal gained from the experiment. Solid line is the same

interface calculated by the heat model.

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Results and Discussion for Electron Beam Melting of Cu and Ta at Different Process Parameters

Using the heat model, numerical experiments are made at different beam powers for investigation of the heat transfer in the refractory metal Ta and in Cu - low-melting-temperature metal. Heat streams are calculated using the obtained linear functions for X and Cp and the results are presented in Table II. The e-beam radius is 10 mm. In both cases of electron beam power (20 kW and 30 kW) significant differences in the energy losses for Ta and Cu can be seen. The most important difference is that there is no contact between the liquid pool and the crucible wall for EBMR of Ta (50 mm ingot diameter), while a contact is observed for Cu (with good thermal conductivity). In the case of Cu, the heat flux through the crucible wall (G2) is the most significant one (54%) for 30 kW beam power and it is also a big part (33%) of the energy losses when the beam power is 20 kW. At e-beam melting of Ta the biggest part of the heat flows is through the ingot bottom G4 due to the small distance between the liquid/solid boundaries and the ingot bottom (the ingot length is 20 mm) despite the relatively low thermal conductivity of tantalum. The radiation losses for the refractory metal Ta are much more than in the case of copper.

Table II. Calculated thermal flows at EBMR of Cu and Ta.

Metal

Cu

' T a

Cu

Ta

Beam power

20 kW

20 kW

30 kW

30 kW

Radiation tosses,W

817

4%

5369

27%

1473

5%

11977

38%

Heat flux through G2,W

6487

33%

0

0%

16359

54%

0

0%

Heat flux through G3.W

696

4%

200

1%

665

2%

378

1%

Heat flux through G4,W

11372

57%

13876

70%

10811

36%

18425

60%

Molten pool volume, %

22.19

10

22.29

14

In order to examine the heat transfer processes and to investigate the influence of the ingot dimensions (ingot length and ingot diameter) during EBMR of Cu, numerical experiments were performed. The ingot length H is 50 mm, 100 mm, 200 mm (the maximal length for our equipment ELIT-60) and the ingot diameter (2R) is 60 mm and 100 mm. The beam power is 20 kW, the beam radius is 10 mm, and the diameter of the feeding material is 50 mm. The casting velocity V is in the range of 1-15 mm/min. The obtained results show that when the ingot length (H) increases the values of the energy losses (radiation and vapor losses) through Gl increase. For H=50 mm these losses are about 1.7 kW (= 8%), for H=100 mm - 3 kW {- 15%), for H=200 mm - 5 kW (= 25%) when the diameter of the metal cast is 60 mm. The heat flow through the interface ingot/puller (G4) increase significantly in the case of shorter ingots because the decrease of the distance between the puller and the reaction surface of the liquid metal pool on the top of the cast block in the crucible. The thermal losses through G4 are: about 16 kW (77 %) for the shortest formed pure ingot (H=50 mm), 7 kW (34 %) for H=100 mm and 2 kW (10-11 %) for H=200 mm for R=30 mm. The results show that the heat stream through the interface cast ingot/crucible wall (G2) decreases significantly with the ingot length decrease and when the ingot radius increases this flow decreases. Considerable part - about 32 % of the energy losses

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passes through this boundary for 60 mm ingot diameter, while for the bigger ingot diameter (100 mm) the liquid pool does not reach the crucible wall at 20 kW. Other numerical experiments were made in order to investigate the influence of process conditions casting velocity (1-15 mm/min) and beam power (20-30 kW) for copper ingots with 60 mm diameter and 100 mm length. The obtained dependencies of the energy losses through the ingot boundaries on the regime conditions are presented in Figs.7-9.

Fig.7. Heat flow through Gl. Fig.8. Heat flow through G2.

Fig. 10. Liquid/ solid boundaries in Cu Fig.9. Heat flow through G4. ingot for three beam powers, V=6

mm/min.

The thermal losses through Gl increase for higher values of casting velocity (Fig.7), because the temperature increase due to the more heat added by the poured molten metal from the melting raw material (Fig.l), causing more vapor and radiation losses. The heat flow through the interface cast ingot/crucible wall (Fig.8), as a part of the total energy losses, decreases when the beam power increases and the flux is about 32% for 20 kW, 30% for 25 kW and 28% for 30 kW. The heat flow through the interface ingot/ vacuum (G3) is about 18%. The values of this stream increase with the casting velocity increase up to 6 mm/min and after that stabilization occurs. Considerable part (about 30% and more) of the thermal losses is through G4 and this heat flow depends strongly on the casting velocity as well as on the beam power (Fig.9). The obtained results show that the heat stream through the bottom of the ingot (G4), for casting velocities up to 6 mm/min, is the main part of the energy losses for the investigated conditions. The lower energy losses through the other ingot boundaries contribute to more efficient EBMR process. These conditions are appropriate to form a flat crystallization front (liquid/solid boundary) which is connected with the quality of the performed pure ingot. Calculated crystallization front shapes in Cu for V=6 mm/min and 20 kW, 25 kW and 30 kW beam powers are presented in Fig. 10. Typical almost flat liquid/solid boundaries are observed. These crystallization front shapes permit formation of dendrite structures parallel to the ingot axis and uniform impurity displacement toward the ingot top surface.

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Conclusion

Using the described heat model, useful for the practice data about geometry of the liquid pool, crystallization front and thermal streams for EBMR of the low-melting-temperature metal Cu (with good thermal conductivity) and the refractory metal Ta at different process conditions is obtained. Variations of the thermal conductivity X and heat capacity Cp vs. the metal temperature are estimated and are taken into account in the quasi-steady-state model. In the case of the metals Cu and Ta with different melting temperatures and thermal parameters (k and Cp), the beam power influence, the heat flows values and liquid pools geometry are very different. The presented results concerning the influence of the ingot dimensions (length and diameter), beam power and casting velocity on the uniformity of the obtained block structure are also very important for the practice. It is shown that flat crystallization front shape for Cu, heated by 20-30 kW, can be obtained for casting velocity up to 6 mm/min. Calculated and experimentally obtained crystallization front forms are compared and a good correspondence is observed. The calculated data about molten pools, heat flows and liquid/solid boundaries can be used for optimization and control of the EBMR process parameters as well as to improve the quality of the produced pure materials.

Acknowledgements

This research was funded by the National Fund for Scientific Research at the Ministry of Education, Youth and Science of Republic of Bulgaria under contracts N° TK01/0073 (DO 02-200/2008) and BIn-5/2009. The authors thank R.Nikolov and D.Tabakov for the technical assistance.

References

1. A. Kalugin, Electron beam melting (Moscow: Metallurgy, 1980) (in Russian). 2. Georgi Mladenov, Electron and Ion Technologies (Sofia: Academic Publ. House, 2009)

(in Bulgarian). 3. I.G. Sharma, N. Krishnamurthy, and A.K. Suri, "Electron beam melting of reactive and

refractory metals and alloys," Proc. of Indo-Bulgarian Workshop on Electron Beam Technologies and Applications, ed. A.K.Das, Mumbai, India, (2004), 42-50.

4. R. Bakish, "Electron Beam Melting 1995 to 2005," Proc. of the 7lh Intern. EBT Conference, Varna, (2003), 233-240.

5. A. Mitchel, T. Wang, "Electron beam melting technology review," In: R. Bakish editor. Proceedings of the Conf. Electron Beam Melting and Refining. Stale of the Art 2000, Reno, Nevada, NJ, USA. (2000), 2-13.

6. S. Sethi, A. Singh, J. Mukherjee, L. Gantayet, "Online measurement of melt-pool depth in EB heated metallic charge for estimation of gap resistance", In: M. Mascarenhas editor. Proc. Int. DAE-BRNS Symp. on Electron Beam Technology and Applications, Mumbai, India, (2005), 343-53.

7. K.W. Westerberg, M.A. Me Clelland, B.A. Finlayson, "Finite element analysis of flow, heat transfre, and free interfaces in an electron-beam vaporization system for metals", Int. J. Numer. Methods in Fluids, 26 (1998), 637-55.

8. G. Gupta, R. Marwah, "Estimation of effective thermal conductivity of hot liquid refractory metal generated by electron beam heating", In: M. Mascarenhas editor. Proc. Int. DAE-BRNS Symp. on Electron Beam Technology and Applications, Mumbai, India, (2005), 371-375.

9. K. Vutova, E. Koleva, G. Mladenov, "Simulation of thermal transfer process in cast ingots

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at electron beam melting and refining", International Review of Mechanical Engineering - special issue on Heat Transfer, 5 (2011), 257-265.

10. K. Vutova, V. Donchev, V. Vassileva, G. Mladenov, "Thermal transfer process through treated metal produced by electron beam melting and refining" (Paper presented at the 1st

CEEC-TAC1, 7-10 September 2011, Craiova, Romania), 265. 11. G. Samsonov, Chemo-physical properties of elements (Kiev: Naukova dumka Publ.

House, 1965) 202-227 (in Russian). 12. S.M. Gurevich, Metallurgy and technology of refractory metals and alloys welding (Kiev:

Naukova dumka Publ. House, 1982) 13 (in Russian).

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INDUSTRIAL USE OF A N E W U L T R A S O U N D S P R A Y F O R C O O L I N

A N D W E T GAS T R E A T M E N T IN THE P Y R O M E T A L L U R G I C A L

PROCESSES

Milorad CIRKOVIC,1 Vlastimir TRUJIC1, Zeljko KAMBEROVIC2

Mining and Metallurgy Institute Bor, Zeleni bulevar 35,Bor, Serbia, e-mail: milorad.cirkovicfaiirmbor.co.rs

2 Faculty of Technology and Metallurgy, Belgrade, Serbia

Keywords: Gas Treatment, Ultrasonic spray, Recycling, Wet dedusting

Abstract

This paper presents the results of a new industrial use the ultrasonic spray for gas cooling from sulphide copper concentrate roasting in the fluo solid reactor and the results of wet gas treatment (washing) in the scrubber plant during treatment the secondary copper bearing raw materials in RTB Bor (Serbia).

The use of a new ultrasonic nozzle for gas cooling from the fluo solid reactor resulted as a necessity because the efficiency of existing installed spray of foreign manufacture was very bad. The use of such spray has caused the formation of large deposits of wet dust on the walls of cooling tower, and their shaking caused the system blockage for dust transport of dust and stoppage the entire plant. Frequent blockage and inability of cleaning resulted into high consumption of nozzles. Construction of a new ultrasonic nozzle was based on the famous Hartman ultrasound generator and it is patented. The possibility of its dismantling (disassembly) has allowed a complete cleaning and reliable operation and the nozzles, currently used, are in a continuous operation over a decade.

New ultrasonic spray has been also successfully applied in the scrubber plant for gas washing during the copper scrap treatment with high zinc content. Separation degree of ZnO in the scrubber from the process gas was 98.6%. The new design of ultrasonic spray and its successful use for cooling and cleaning the metallurgical gases will be applied in modernization the Copper Smelter in RTB Bor.

Introduction

When it is discussed on ultrasound and ultrasound technique, the problem is usually studied from the perspective of specific applications in various scientific, medical, technical, technological and other fields. Engineers designers and scientists are faced with a difficult task when they approach to development of new, targeted ultrasound devicesnthat have to perform some technological operation and to serve in certain measuring - control systems.

Within the modernization of system for treatment the gases from the process of roasting the copper concentrates in fluo-solid reactor of the Copper Smelter in RTB Bor, in 1982 a new plant was installed, of foreign manufacturer, which is a part of cooling tower and electrostatic precipitator. In the cooling tower, a partial cooling and desting of gases is carried out (temperature degradation of 550oC to 360°C) prior to their final dedusting in the electrostatic

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precipitator. Three ultrasound sprays were installed in the cooling tower for gas cooling. Shortly after the start of this plant in the cooling tower, the great problems occurred that

worsened the operation of this unit. On the walls of tower, the large stickersof wet dust were formed that fell out and blocked a transport system on the bottom in which a deadlock of technological line for roasting and smelting the copper charge came. In addition to this, water often occurred on the bottom of cooling tower that did not evaporate during cooling process, and the second part non-evaporated water caused formation of wet dust that was deposited on electrode system and blocked the work of electro filter. The result of occurred problems are frequent delays in production and deterioration of production results in copper production.

The cause of given technological problems is poor operation of installed ultrasonic sprays, i.e. their non-functional design has caused the operation blockage. Impurities from water and carbonate sedimentation in the water holes blocked the certain openings and it caused the diversion of water spray and wetting the tower walls. In this case, the effect of ultrasound atomization of water could not be achieved and therefore instead of submicron water droplets, more larger drops appeared that falls out on the bottom of cooling tower and partly went into electrostatic filter.

In order to eliminate these problems and improve the operation of system for gas treatment of fluosolid reactor, it was necessary to construct a new functional ultrasonic atomizer for water. After extensive scientific researches, a new ultrasonic spray was designed, based on known the Hartman ultrasound generator, which was installed in the cooling towers of both technological lines for roasting and melting in 1994 in the Copper Smelter of RTB Bor, which is now in operation. The new ultrasonic spray, which is now in operation, is patented, [4].

Ultrasound sprays, due to their reliable operation and extensive experience in their exploitation, were applied in scrubber installation in the copper scrap treatment with high zinc content. In this case, a wet treatment (washing) of gas was achieved that contained large amount of ZnO. Unlike the conventional sprays, the ultrasound sprays showed better efficiency in treatment the metallurgical gases due to a phenomenon of coagulation the solid particles in gases in the ultrasound field. These experimental studies have shown that consumption of water for gas washing using the ultrasonic sprays was significantly lower compared to the application of other sprays.

Reliability and efficiency of use these new ultrasonic spray will be significant in next period in which the Copper Smelter in RTB Bor will be modernized. The plan of modernization the Copper Smelter envisages installation a new autogenous smelting technology, and also the need will be for cooling the metallurgical gases, where a domestic ultrasonic spray has a considerable advantage.

Starting problem of cooling the fluosolid reactor gases of RTB Bor

The problems of gas cooling and dedusting of the process of roasting the sulphide copper concentrates in the Copper Smelter in RTB Bor was expressed after installing a new system for process gas treatment during reconstruction both technological lines in the Copper Smelter in 1982. Both technological lines of roasting and smelting have one plant for reactor gas treatment consisting of a cooling tower and electrostatic precipitator. The first stage in treatment of reactor gases is cooling and partly dedusting in the cooling tower. Figure 1, schematically shows the cooling tower with gas characteristics and cooling fluids.

Gas cooling gas in a cooling tower is realized by water that is atomized using three ultrasonic sprays. The original spray, i.e. its atomizer for water, has a complex structure andconsists of five elements. Using this spray in the operating conditions can cause the additional difficulties. These sprays are in the group of sprays with two fluids out of which one

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is compressed air and serves as a working fluid for ultrasound production and water atomisation. At the start of plant for treatment the reactor gases, and especially cooling tower, some

problems appeared, caused by faulty operation of cooling water sprays. The main problem was that the spray nozzles were often clogged with impurities from water, and cleaning of nozzles could not be fully realized. In the cases, the loss of jet geometry of atomized water also led to damage of resonator and frequent replacement of this spray element.

Figure 1. Review the cooling tower with characteristics of gas and cooling fluids

Constructively, the nozzle is constructed as unseparable circuit so that the impurities deposited in the openings for water in the outlet area and blocked the water flow. Attempts to clean the nozzle caused damages the channels for water. These damages caused an incomplete atomization of cooling water and its incomplete evaporation, diversion of jet and spraying the walls of cooling tower, and often water appeared in the output conveyor. In this case, large stickers of dust are formed on the walls of tower which often fell out and block the work of transporters on the bottom of tower and the unit must be stopped for cleaning. Frequent dismantling for cleaning sprays caused damages in all its elements and thus large consumption of nozzles and other parts of sprays.

High consumption of ultrasound spray parts and often delays in production resulted into deterioration in the economics of copper production. Such situation was unsustainable and the solution was necessary to find in the substitution of foreign with domestic sprays that will be better and solve the existing problems.

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New ultrasonic spray for gas cooling of roasting process the copper concentrates in the fluosolid reactor in RTB Bor

Theoretical basis

Construction the new ultrasonic spray for cooling the roasting gases in the fluosolid reactor of the Copper Smelter in Bor and substitution the original imported ultrasonic spray was an imperative in order to reduce delays in the production and elimination of import a large quantities of supersonic sprays parts.

After extensive scientific researches, a new ultrasonic spray was designed, based on known the Hartman ultrasound generator. This is the type of generator which uses the kinetic energy of compressed gas for ultrasound production. The working principle of the Hartman gas generator is based on appearance the leakage of gases, higher pressure through a narrow slot of nozzle, Figure 2, [2].

Figure 2. Theoretical basis of the Hartman ultrasound generator

When a jet of compressed air (J) leaks out from the nozzle (N) then a field of variable pressure appears in it, Figure la, lb. This pressure decreases with moving away from the nozzle opening, Figure 2c, 2d. Distances ai,bi,a2,b2, in which the points of minimum and maximum pressure appear, depend on the initial pressure (p) and hole of nozzle (do). The fields of pressure between points ai-bi and a2-b2, Figure 2c, are called the fields of instability and they have an important role in the creation of ultrasound. If the resonator, cylindrical cavity, Figure 2b, is set in the field of instability, there will be its alternately filling and emptying with a frequency

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determined by dimensions of the cavity, i.e. resonator. In this way,a high-intensity ultrasound is created.

New ultrasound spray in the Copper Smelter RTB Bor and the results of Its use

The ultrasound spray is designed on the exposed base that was installed on the cooling tower for cooling gases of the roasting process of sulphide copper concentrates in the Copper Smelter in Bor, Figures 3 and 4.

Figure 3. Ultrasound spray Figure 4. Constituonal parts of a new ultrasound spray a) outer nozzle sleeve, b) inner nozzle section, c) resonator

Nozzle is designed as a detachable fit and it consists of the outer sleeve, Figure 4a, and inner section, Figure 4b, which are joined together by thread. This construction of nozzle is a high technical improvement because it allows rapid assembly and disassembly. Possibility of nozzle disassembly allows its complete and easy cleaning which is essential difference compared to the previous import nozzle.

Resonator is an important part of ultrasound spray and it is a generator of ultrasound. It is made of solid material and has much better characteristics, and life time is five times longer than imported one. During the manufacturing process in the Copper Smelter in RTB Bor, the plant parameters were controlled for gas treatment of fluosolid reactor and visual control of inner lining the cooling tower. It is very important that the designed parameters of cooling gas are fully realized and that substitution was successfully realized.

Exploitation of a new ultrasonic spray lasts continuously for almost two decades. In that period, any nozzle was not replaced but the nozzles that were first built are inexploitation. Resonators were changed several times during such long period of exploitation.

Application of ultrasound sprays in the scrubber plant for gas washing

New ultrasonic spray has been also successfully applied in the scrubber plant for gas washing during the copper scrap treatment with high zinc content. Industrial scale testings were performed in the plant-site of RTB Bor. For the aim of successful realization of experiments, and

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before all maximum safety and removal of risks from toxic pollution the working environment, the plant for gas dedusting from the process of brass scrap treatment was constructed in the site of Copper Foundry in RTB Bor and presented in Figure 5, [1,2,3,5,6].

Figure 5. Schematic representation of the plant where industrial scale experimental investigations were carried out in the plant of Bor Company 1-smelting furnace; 2 - cooling chamber; 3 - scruber; 4 -fan for gas drawing; 5 - stack; 6 - ultrasound sprays.

Scruber of steel sheets was constructed for the needs of wet dedusting. It consists of cylindric and cone part, supply and discharge gasline, [2]. Cylindrical part of scruber has length of 2000 mm and diameter of 1200 mm, and cone part length is 2500 mm. Scruber is supported by four steel holders, and is installed such as the lowest part is at height of 1100 mm over basement. Container for ZnO is situated in this area. Three experiments of brass scrap remelting were carried out, and temperature of exhausted gases from furnace was in limits 1000-1200°C.

Two ultrasound sprays were installed into cooling chamber, and three sprays for water atomization were installed into scruber.

Table 1, presents quantities of metal scrap as well as copper and zinc quantities in each charge, also zinc oxide quantities that have to be seprated in scruber installation.

Table 1. Content of charges in remelting process of brass scrap No.of operation

1 2 3

Scrap quantity (kg) 4301 5065 5105

Cu (kg)

3129.4 3685.3 3714.4

Zn (kg)

1153.5 1358.4 1369.4

others (kg) 18.1 21.3 21.2

ZnO (kg) 1435.9 1658.4 1697.8

It is seen from Table 1, that zinc oxide quantity separated from process, is significant, and that its separation from process gases presents a great problem. Separation of this compound in the process of wet gas treatment is an obligation, because discharge and very small quantities of toxic Zn would jeopardize safety of working environment.

Table 2 and Figure 6, presents a change of chemical content of molten metal during smelting operation and zinc removal during experiment No.l. During experimental smelting, samples of

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molten metal were taken, and copper and zinc content was analyzed depending on operation time.

Table 2. Change of chemical composition of molten metal in experiment No I

Sample mark

S- l S -2 S - 3 S-4 S - 5 O - l 0 - 2 0 - 3 0 - 4 0 - 5 0 - 6 0 - 7 0 - 8 0 - 9 O-10 R - l R - 2 R - 3 R - 4

Sampling interval min

118 189 226 281 306 351 481 531 563 618 656 701 736 756 779 799 824 826 858

Cu % 73,98 74,22 76,00 77,63 77,51 80,63 84,90 86,21 88,56 94,33 93,99 95,78 98,43 99,02 99,14 98,60 98,76 98,88 99,03

Zn % 25,75 24,86 23,25 21,99 21,66 18,53 14,34 13,01 10,41 4,95 5,26 3,05 1,03 0,50 0,34 0,92 0,77 0,70 0,45

Figure 6. Change of copper and zinc of molten metal in exp. No 1 depending on operation time

S-smelting; O-oxidation; R-reduction

During tretment process of brass scrap, ZnO content was controlled in inlet gases. Quantity of zinc oxide in outlet gases was 3.069 kg/h at measuring point No.2 from scruber. Total quantity of ZnO, exhausted with process gas into atmosphere over stack was 19.9 kg, what presents minimum quantity for those conditions.

It means that efficiency of such simplified system for gas treatment was 98.6%, and total water flow through sprays in this process was 4920 m3/h.

Similar results were obtained in other two experiments, and it could be said that the given task was successfully realized.

Figure 7, presents a photo of scruber installation during experiments, and Figure 8, presents a photo of zinc oxide discharhe in a form of pulp from cone part of scruber.

Figure 7. Scruber installation Figure 8. Zinc oxide separation

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Conclusion

Based on the previous industrial use of a new ultrasonic spray for gas cooling of roasting process the copper concentrate in the fluosolid reactor in the Copper Smelter in RTB Bor and control of the site designed parameters, it can be concluded the following:

- New local spray successfully replaced the imported spray, - Possibility of nozzle disassembling enables its fully cleaning what extends its

exploitation life time, - All negative effects of wotk the original spray for gas cooling are removed and now gas

treatment system works completely stable, - Delays and negative effects in the copper production are significantly reduced and the

economic effect was improved to the amount of 0.75 to 0.8% per annum and with increase of xoncwnträte treatment and copper production, reducing downtime, etc.,

- In the upcoming modernization of the Copper Smelter in RTB Bor and installation of modern autogenous technology of smelting the copper concentrates, a new ultrasound spray for cooling of metallurgical gases will have an important role having regard that it is the domestic construction and tested in practice.

The use of new ultrasonic sprays was successfully implemented in the experimental researches in the scrubber installation for wet cleaning of gases in the process of copper scrap treatment with the increased zinc content. Based on the results of three experimental smeltings the brass scrap, the following could be concluded :

- Defined system, that is constructed plant, was in function without any problems during experimental testings, and that is pointed out by degree of ZnO separation of 98.6 %,

- Sprays with ultrasound atomization of cooling water and wet gas dedusting, used for the first time in such system, have shown as very efficient and favourable,

- Working area during experiments was safe for work, i.e. there is no ZnO emission out of the plant,

- Industrial experimental investigations were developed undisturbedly.

References

1. Hudjakov I.F., Doroäkeviö A.P., Karelov S.V., Metallurgija vtoriênih tjazolih cvetnih metallov, Metallurgija, Moskva,1987.

2. Alesïna V.M., Pileulavlivanie v metallurgii, Metallurgija, Moskva, 1984. 3. Kovacevic Dj., Ultrasound and its Use in Industry, Technical Book, Belgrade, 1958.(in Serbin) 4. Cirkovic M., Patent No. 866 MP (in Serbian) 5. Kocovski B., Copper and Copper Alloys, Copper Magazine, Bor,1991.(in Serbian) 6. Babadzan A.A., Pirometallurgiceskaja selekcija, Metallurgija, Moskva, 1968.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DEVELOPMENT OF 3D POROUS NICKEL ELECTRODES FOR HYDROGEN PRODUCTION

Valentin Pérez-Herranz', Isaac Herrâiz-Cardona1, Emma Ortega1, José Garcia-Anton'

'lEC Group, Depto. Ingenieria Quimica y Nuclear, Universidad Politécnica de Valencia Camino de Vera s/n, 46022 Valencia, Spain.

Keywords: 3D Ni Electrodeposits, Dynamic Template, Surface Roughness Factor, Electrochemical Impedance Spectroscopy

Abstract

The production of hydrogen is of particular interest because hydrogen is considered as a clean fuel. Hydrogen production via electrolysis of water from alkaline aqueous electrolytes is a well-established conventional technology. However, the high energy consumption retrains its large-scale application at present. This work focusses on the development of nickel electrodes with excessive surface areas. The construction of 3D porous structures was achieved by means of a hydrogen bubble dynamic template, prepared from electrochemical deposition processes. Cu was electrodeposited and grew within the interstitial spaces between the hydrogen bubbles to form a macroporous film of Cu nanoparticles on the substrate. Subsequently, a Ni layer was electrodeposited on the Cu template. The hydrogen evolution reaction was evaluated in alkaline media by means of steady-state polarization curves and electrochemical impedance spectroscopy. The developed electrodes exhibited a significantly reduced overpotential for the HER when comparing to the commercial smooth Ni electrode.

Introduction

Hydrogen is considered an ideal energy carrier that can be an alternative to fossil fuels. It is a clean and fully recyclable substance with a practically unlimited supply, and has all the criteria considered for an alternative energy source. The electrochemical production of hydrogen by alkaline water electrolysis is one of the most promising methods with great potential of using renewable energy sources [1]. However, the high energy consumption of alkaline water electrolysers retrains its large-scale application at present. Among the electrode materials, nickel and its alloys show a high initial electrocatalytic activity toward the HER, that can be enlarged by increasing the real surface area and/or the intrinsic activity of the electrode material [2].

The increase of the real surface area can be achieved by several methods: depositing Ni together with an active metal like Al or Zn [3-7] followed by the dissolution of the secondary component (Raney type electrodes); electrodeposition of Ni at large current densities [8-9], electrodeposition of Ni on metallic opals (made of silica or polystyrene) with proper porosities and layer/thickness, followed by a selective removal of the opal [10]. As a result, a porous, three-dimensional (3D) structure is obtained, characterized by a high surface roughness factor, Rf.

In this work, the electrochemical preparation of 3D porous Ni electrodes for HER is reported. These electrodes were synthesized by nickel electrodeposition on copper foams obtained from hydrogen bubbles dynamic templates. The electrocatalytic performance of the developed electrodes for HER was evaluated in 30 wt.% KOH solution by using polarization curves and electrochemical impedance spectroscopy (EIS) techniques.

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Experimental

The metallic coatings were deposited on an AISI 304 stainless steel substrate, embedded in Teflon, leaving a cross-sectional available area of 0.5 cm2. Before the electrodeposition experiments, the stainless steel substrate was mechanically polished with emery paper down to 4000 grit, next it was degreased for 1 min with 25 wt.% NaOH at 90 °C, immersed in HC1 18 wt.% during 1 min and anodically treated in 70 wt.% H2S04 at 1080 A m"2 for 3 minutes. Then, the substrate surface was struck at 268 A m"2 in a Wood's nickel solution (240 g L"1 NiCl2, 120 mL L"' HC1) for 5 min, in order to produce a thin, adherent deposit of nickel which serves as a base for the subsequent electrodeposition. Between each treatment the electrode was rinsed with distilled water.

The formation of electroactive coatings was done by electrodeposition. First of all, it was constructed a 3D porous copper structure by means of electrodeposition at high current densities in an acidic bath and at room temperature. Afterwards, the Cu macroporous layer acted as a template for the nickel deposition. Nickel was electrodeposited from a modified Watts bath at 50°C. The bath composition and deposition conditions of the tested electrodes are listed in Table 1. The structures, morphologies and compositions of the developed materials were examined by means of a JEOL JSM-3600 scanning electron microscope (SEM) coupled with an Energy Dispersive X-Ray (EDX) Spectrometer, and an OLYMPUS LEXT OLS3100-USS confocal laser scanning microscope. Electrodepositions were carried out in a one-compartment cell. The counter electrode was a large-area platinum electrode. The reference electrode was an Ag-AgCl electrode with 3 M KC1 solution.

Table 1. Bath compositions and operating conditions used in the electrodeposition of electrocatalytic coatings on an AISI 304 stainless steel substrate.

The developed electrodes were characterized by means of polarization curves and electrochemical impedance spectroscopy (EIS). All these tests were performed in oxygen free 30 wt.% KOH solutions. Polarization curves were potentiodynamically recorded at a scan rate of 1 mV s"1, and at six different temperatures: 30, 40, 50, 60, 70 and 80 °C. EIS measurements were performed at different cathodic overpotentials, and at the following temperatures: 30, 50, and 80 °C. The measurements were made in the frequency range of 10 kHz to 3 mHz. Ten frequencies per decade were scanned using a sinusoidal signal of 10 mV peak-to-peak. All the electrochemical experiments were performed using an AUTOLAB PGSTAT302N potentiostat/galvanostat.

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Results and discussion.

Fig. 1 shows the confocal laser scanning microscopy images of the porous Cu-Ni alloys synthesized according to the electrochemical deposition conditions reported in both Table 1. The deposited films are characterized by a 3D foam structure. As depicted in Fig. 1, the macropore size of the foam increases with both the time of deposition and the decrease in the applied current density. The walls of the copper structure obtained by this method are composed of numerous ramified deposits which constitute some nanoparticles (~300 nm), finger-like structure. Fig. 2 shows the effect of nickel deposition on the microstructure of the 3D porous copper templates. As depicted in Fig. 2, as the time of Ni electrodeposition increases, the microstructure of the coatings changes from the finger-like to the cauliflower structure. It is important to control the nickel electrodeposition conditions, due to the fact that a thicker nickel layer blocks the macroporous structure, as it is observed in Fig. 1 for the ED-1 electrocatalyst.

Fig. 1. Confocal laser scanning micrographs of the different obtained electrodes: a. ED-1, b. ED-2, c. ED-3, and d. ED-4.

Fig. 2. SEM images of the macroporous 3D Ni electrodes after different Ni electrodeposition time: a. 15 min; b. 30 min; c. 60 min.

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The catalytic activity of the prepared layer was evaluated by means of Tafel linear polarization measurements in 30 wt.% KOH solution. Fig. 3 shows a set of Tafel curves recorded at 50 °C on the catalyst coatings investigated. A curve performed on commercial smooth Ni electrode was also included to compare the obtained results. The curves were corrected with respect to the reversible HER potential at the given conditions and for the/ft-drop. The information obtained from the Tafel polarization data demonstrate that all the investigated catalysts are very active for the HER, showing a higher catalytic activity than the smooth Ni electrode. Between all the tested 3D nickel eiectrocatalysts, ED-1 has the lowest catalytic activity for HER, which is attributed to a more intense nickel electrodeposition that considerably fills the macropores, generating a smoother layer and, consequently, a decrease in the active surface area. On the other hand, between ED-2 and ED-3 eiectrocatalysts it can be appreciated two different regions: at overpotentials less cathodic than -150 mV, ED-3 electrode manifests a better catalytic activity than ED-2 coating, whereas this trend is inverted at more cathodic potentials. The reported behaviour is directly connected with the superficial morphology. The macropores of ED-3 electrocatalyst are blocked by gas bubbles when HER is vigorous, due to their smaller size, leading to low utilization of the whole surface. In contrast, the higher pore size of ED-2 catalyst facilitates the fast transport of electroactive gas through the porous electrode, increasing the hydrogen production for a same applied overpotential. Therefore, from Fig. 3 it is clear that the higher pore size of the catalyst surface the higher electrode activity, being ED-4 the best overall catalyst.

Fig. 3. Linear Tafel polarization curves recorded on the different developed electrodes in 30 wt.% KOH solution at 50 °C.

To ensure a complete characterization of the electrode/electrolyte interface and corresponding processes, EIS measurements were made at different selected overpotentials. Fig. 4.a and 4.b show examples of EIS spectra recorded on the ED-2, and ED-4 eiectrocatalysts, respectively. The EIS spectra recorded on the ED-2 coating reveal the presence of two strongly overlapped semicircles (i.e. two different time constants). This behaviour was also observed for both ED-1 and ED-3 electrodes. With respect to the impedance spectra of the ED-4 electrode, it is shown that the two semicircles are clearly differentiated. In both cases, at the higher cathodic overpotential applied, HER is so vigorous that hydrogen bubbles cause too much interference, good impedance spectra cannot be obtained and only one deformed semicircle appears. From Fig. 4.a and 4.b it is clear that the diameter of both semicircles considerably diminishes with the cathodic overpotential for all the investigated electrodes. The same phenomenon was evidenced

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with the increase in the temperature, indicating that both time constants are related to the electrode kinetics [11].

To model the experimental data of the impedance response characterized by one semicircle in the Nyquist plot (i.e. recorded at highest overpotentials) it has been used the one time constant (IT) electric equivalent circuit (EEC) model. This model is the classical Randies EEC in which the double layer capacitance was replaced by a constant phase angle element (CPE), see Fig. 5.a [12]. CPE is defined in impedance representation as:

where Q is the CPE constant, m is the angular frequency (in rad/s), i2 = -1 is the imaginary number, and n is the CPE exponent.

With respect to the impedance response characterized by two semicircles in the Nyquist plot, according to the behaviour of the high frequency semicircle with both overpotential and temperature, the two-time constant parallel (2TP) EEC model shown in Fig. 5.b has been used to fit the EIS response of the catalysts investigated [4, 11, 12], The 2TP model reflects the response of a HER system characterized by two semicircles (i.e. two time constants): the high frequency (HF) semicircle, ti (CPEi, Ri), related to the charge transfer kinetics, and the low frequency (LF) semicircle, 15 (CPE2, R2), related to the hydrogen adsorption [4, 11, 13, 14].

Fig. 4. Nyquist representation of the impedance data obtained in 30 wt.% KOH solution at 30 °C for the electrodes: a. ED-2, and b. ED-4. Symbols are the experimental points and solid lines are

modelled data.

Fig. 5. EEC models used to explain the EIS response of the HER on the developed electrocatalysts: (a) one-time constant (IT), and (b) two-time constant parallel model (2TP).

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Table 2 shows the best-fit estimates of the different EEC parameters obtained from the impedance measurements of both ED-2 and ED-4 electrodes at different overpotentials and temperatures. The average double layer capacitances, C„ for the catalytic coatings were determined using the relation suggested by Brug et al. [15]:

C( =[&/(* / '+* ." ' )0 -1 0 ]"* (2)

Table 2. EEC parameters obtained by fitting EIS experimental spectra recorded at various overpotentials an temperatures in 30 wt.% KOH solution on the investigated electrocatalytic coatings.

According to the EEC parameter values presented in Table 2, the two time constants, T\ (CPE], Ri) and T2 {CPE2, R2), decrease with the cathodic overpotential and the temperature, indicating that both semicircles are related to the kinetics of the process [11]. As C; and the Rj decrease with cathodic overpotential, the first time constant, n or HF time constant, is related to the HER charge-transfer kinetics. On the other hand, it can be distinguished two different behaviours of the second time constant, T2 or LF time constant, for the tested electrodes. With an increase in the cathodic overpotential, the value of C2 increases, while the value of R2 decreases for the ED-4

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cathode. This is a typical behaviour related to the response of hydrogen adsorbed on an electrode surface [4, 11, 13, 14]. In contrast, in the case of ED-1, ED-2, and ED-3 electrocatalysts, the C2

and the R2 rapidly decrease with overpotential.

Besides the information on the kinetics of the HER, EIS results can be also used to estimate the real surface area of electrocatalytic coatings. Considering a value of 20 uF cm"2 for the double layer capacitance, CM, of a smooth nickel surface [16], the real active surface area, in terms of surface roughness factor (Rf), may be estimated by comparing the CM related to the HER charge-tranfer kinetics (Ci) of porous/rough and smooth electrodes [17], The plots of the electrode Rfas function of the HER overpotential are displayed in Fig. 6. As it is clear from Fig. 6, the values of Rf decrease when increasing the cathodic potential. This indicates that a fraction of the inner surface of the electrode is blocked during HER due to gas bubbles shielding, and hence not electrochemically accessed by the electrolyte [18],

Fig. 6. Surface roughness factor, Rf, as a function of the overpotential for the developed electrodes in 30 wt.%. KOH solution at 50 °C.

Conclusions.

Three-dimensional foam Ni structures were successfully created by a double-template electrochemical procedure. The hydrogen evolution reaction (HER) on these electrodes was assessed by polarization curves and electrochemical impedance spectroscopy (EIS). This research allowed us to enhance that:

1. The investigated catalysts manifest greater apparent activity towards HER in comparison with commercial smooth Ni electrode.

2. Copper opals obtained at lower current densities and higher deposition times (ED-4) lead to 3D Ni catalysts with a higher pore size. The higher the pore sizes of the macroporous catalyst, the higher the electrode activity.

3. From both linear Tafel polarization curves and EIS it is derived that the increase in the catalytic activity of the developed electrodes is mainly attributed to an increase in the electrochemical active surface area.

Acknowledgements.

Isaac Herraiz-Cardona is grateful to the Ministerio de Educaciön (Spain) for a postgraduate grant (Ref. AP2007-03737). This work was supported by Generalität Valenciana (PROMETEO/2010/023) and Universidad Politécnica de Valencia (PAID-06-10-2227).

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[7] M. Okido, J.K. Depo,and G.A. Capuano, "The mechanism of hydrogen evolution reaction on a modified Raney-Nickel composite-coated electrode by AC impedance," Journal of the Electrochemical Society, 140(1993), 127-133.

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[10] Y.J. Huang et al., "Ni inverse opals for water electrolysis in an alkaline electrolyte," Journal of the Electrochemical Society, 157(2010), 18-22.

[11] L. Birry, and A. Lasia, "Studies of the hydrogen evolution reaction on Raney nickel-molybdenum electrodes," Journal of Applied Electrochemistry, 34 (2004), 735-749.

[12] L.L. Chen, and A. Lasia, "Study of the kinetics of hydrogen evolution reaction on Nickel-Zinc alloy electrodes," Journal of the Electrochemical Society, 138 (1991), 3321-3328.

[13] E. Navarro-F lores, Z.W. Chong, and S. Omanovic, "Characterization of Ni, NiMo, NiW and NiFe electroactive coatings as electrocatalysts for hydrogen evolution in an acidic medium," Journal of Molecular Catalysis A: Chemical, 226 (2005), 179-197.

[14] G.J. Brug et al., "The analysis of electrode impedances complicated by the presence of a constant phase element," Journal of Eledroanalytical Chemistry, 176 (1984), 275-295.

[15] E.B. Castro et al., "An electrochemical impedance study on the kinetics and mechanism of the hydrogen evolution reaction on nickel molybdenite electrodes," Electrochimica Ada, 42 (1997), 951-959.

[16] A. Rami, and A. Lasia, "Kinetics of hydrogen evolution on Ni-Al Alloy electrodes," Journal of Applied Electrochemistry, 22 (1992), 376-382.

[17] S. Trasatti, "Electrocatalysis: understanding the success of DSA," Electrochimica Ada, 45 (2000), 2377-2385.

[18] B. Pierozynski, and L. Smoczynski, "Kinetics of hydrogen evolution reaction at Nickel-Coated carbon fiber materials in 0.5 M H2S04 and 0.1 M NaOH solutions," Journal of the Electrochemical Society, 156 (2009), B1045-B1050.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ELECTROCHEMICAL RECOVERY OF ZINC PRESENT IN THE SPENT PICKLING BATHS COMING FROM HOT DIP GALVANIZING

PROCESSES

Valentin Pérez-Herranz1, Jordi Carrillo-Abad1, Montserrat Garcia-Gabaldon', Emma Ortega'

'iEC Group, Depto. Ingenieria Quimica y Nuclear, Universidad Politécnica de Valencia Camino de Vera s/n, 46022 Valencia, Spain.

Keywords: Electrochemical deposition, Iron, Hydrogen evolution reaction, Pickling solutions, Zinc electrodeposition.

Abstract

Hot dip galvanizing processes offer a simple and effective method for corrosion protection of steel. In this process, during the pickling step, HC1 reacts with iron and iron oxides. Spent pickling baths contain hydrochloric acid, ZnCl2 and FeCl2 as principal compounds. Due to the inadequate existing techniques to treat the spent pickling solutions, the decrease of natural reserves of non-ferrous metals and the requirement of environmental protection, in the present work, the cathodic electrodeposition of zinc present in the spent pickling baths coming from hot dip galvanizing industries is studied. As the electrode potential was shifted towards more negative values, the fractional conversion increased. Simultaneously, the specific energy consumption decreased initially due to the increase in the zinc conversion rate but decreased for the most cathodic potential value due to hydrogen evolution reaction. Even though iron deposition does not take place for any experimental condition under study.

Introduction

Hot dip galvanizing processes offer a simple and effective way for corrosion protection of steel parts. This process consists of the following steps: alkaline or acidic degreasing, rinsing with water, pickling with dilute hydrochloric or sulphuric acid, rinsing with water, fluxing in aqueous ZnCl2/NH4Cl baths, drying and dipping into molten Zn at temperatures of about 450°C for a defined period [1].

As steel is introduced into the pickling bath, HC1 reacts with iron and iron oxides to form ferrous chloride in solution. The same bath is used until the acid concentration falls to give unacceptable pickling times. Spent pickling baths contains hydrochloric acid, ZnCl2 and FeCl2 as principal compounds in the range of 80-150 g/1 FeCl2, 5-150 g/1 ZnCl2 and 10-80 g/1 HC1 [2].

Currently, many techniques exist to eliminate zinc and iron from wasted pickling solutions coming from hot dip galvanizing processes, such as precipitation-filtration, diffusion dialysis, electrodyalisis and solvent extraction. [3]. However, these techniques do not allow any recovery of the products. Then, due to the inadequate existing techniques to treat the spent pickling solutions, the decrease of natural reserves of non-ferrous metals and the requirement of environmental protection, zinc electrowinning from spent pickling solutions appears an interesting alternative [4, 5]. In the present work, the cathodic electrodeposition of zinc present in the spent pickling baths coming from hot dip galvanizing industries is studied under potentiostatic operation.

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Experimental

The electrochemical reactor used in this work consisted of a Pyrex glass of 100 ml with two graphite electrodes acting as working and counter electrodes and a standard Ag/AgCl satured KC1 electrode acting as reference electrode. Both anode and cathode were made of two cylindrical graphite bars with an effective area of 14.15cm2. The potential values to be applied in the electrolysis experiments were chosen from previous voltammetric experiments and ranged between -IV and -1.75V. All the experiments were performed using an Autolab PGSTAT20 potentiostat/galvanostat.

Samples were taken from the electrochemical reactor every 30 min, and potential, current, cell voltage, pH and metal concentrations were recorded during the electrolysis. The determination of zinc was performed by atomic absorption spectrophotometry (AAS) on a Perkin-Elmer model AAnalyst 100 atomic absorption spectrophotometer using a Zn hollow cathode lamp at 213.9 nm wavelength, 0.7 nm spectral bandwidth and an operating current of 5 mA. To measure iron concentration it was used the same equipment changing the Zn hollow lamp for a Fe hollow lamp, the wavelength used was 248.3 nm, the applied operating current was 5 mA and the spectral bandwidth was 0.2 nm.

All solutions were prepared using analytical grade reagents and distilled water. Electrolytes containing ZnCb and/or FeCb in HC1 were used in a concentration range similar to that present in the spent pickling baths. The zinc and iron concentrations were respectively 0.055 M and 0.035 M, which are in accordance with the concentration values of these metal ions present in the diluted pickling-wasted baths. The HC1 concentration was 0.1 M. All the experiments were performed at room temperature.

Results and discussion

Initially, the electrolysis of synthetic solutions in absence of iron has been studied, and, afterwards the iron effect over zinc electrodeposition has been evaluated. Fig. 1 shows a plot of the evolution of zinc concentration as a function of time at different electrode potentials for an electrolyte composed of 0.055M ZnCk and 0.1M HC1. When the electrode potential is -1.125V the concentration of zinc in solution remains practically constant, because this potential value is far from the potential value corresponding to the electrochemical reduction of Zn2+ in the conditions under study [6]. For more cathodic applied potentials, the concentration profile decays exponentially with time as zinc is depleted from solution, this behaviour being characteristic of a batch electrochemical reactor. On the other hand, the more negative the applied potential is, the higher the velocity of the zinc reduction reaction, and consequently, the amount of zinc in solution is lower. At the most cathodic potential value of-1.5V and for a high electrolysis time, an increment in the zinc concentration value is observed. This fact is related to zinc redissolution at high local pH values around working electrode interface [7, 8].

Fig. 2 represents the evolution of current versus time for the same experimental conditions as those presented in Fig. 1. When the applied potential is more cathodic than -1.125V, the current curve shows a maximum value, which is the consequence of two opposite effects. Current sharply rises during the initial period as zinc nucleates on the graphite electrode, causing an increase in roughness, and thus, in total current [8, 9], For longer times, zinc deposition begins to be diffusion controlled, and as zinc ion is depleted, current slowly decreases, this factor being

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predominant at longer times of the experiment. In the case of a potential value of-1.125 V, since Zn deposition is almost negligible the maximum in current is not observed. If the evolution of current is compared for every potential value, it is observed that the more negative the applied potential is the higher the cathodic current value is, due to the greater rate of zinc electrodeposition.

Fig. 1. Evolution of zinc concentration vs. time as a function of the applied potential. [ZnCl2]o = 0.055M, [HCl]o = 0.1M.

Fig. 2. Current intensity vs. time as a function of the applied potential. [ZnCl2]o = 0.055M, [HC1]0 = 0.1 M

In order to rationalize the performance of a given electrochemical reactor or to compare between the performance of electrochemical reactors and other processes, it is essential to consider several figures of merit, such as current efficiency, <|>, the mean value of the space-time yield, r|, and the specific energy consumption, Es, which are calculated with the following equations [10]:

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m=V(C»-C«)) ( g . / . , , - , ) ( 2 )

\u(t)-l{t)-dt , .

Figs. 3-5 show, respectively, the variation of the current efficiency, the space-time yield and the specific energy consumed at different electrode potential values and for an electrolyte composed of 0.055M ZnCb and 0.1M HC1. At die lowest cathodic potential value, the current efficiency falls due to the zinc reduction reaction slowness and the greater importance of parallel reactions [11] (Fig. 3). The current efficiency increases as the electrode potential is made more negative because of the approach to the potential value of the zinc reduction reaction, but at -1.5V hydrogen evolution is more important than zinc deposition and, consequently, this fact causes a decrease of <() in relation to the values obtained at -1.25V and -1.375V. For all the applied potentials under study, <|> presents high values during the initial steps but it drops at longer times due to the decrease in zinc concentration and the increase of hydrogen evolution reaction. Since the HC1 concentration in solution is quite high (0.1M), hydrogen evolution appears from the early stages of all experiments [6], and the maximum value of the current efficiency obtained is close to 45%.

For the lowest cathodic potential value of-1.125V, the space-time yield is negligible and remains practically constant due to the slower rate of the zinc reduction reaction (Fig. 4). For electrode potentials more negative than -1.125V, n presents high values during the initial period and it may be attributed to the nucleation of zinc onto the graphite electrode. The space-time yield slow decrease at longer times is related to the removal of zinc from solution. For a given time value, the space-time yield increases when the cathodic potential becomes more negative, because of the turbulence promoting action of hydrogen evolution, the roughness increase of the electrode surface and the higher velocity of the electrochemical reactions [12].

As shown in Fig. 5, when the electrode potential is -1.125V, the specific energy consumption is quite high because the zinc reduction reaction rate is still low, but when the applied potential is made more cathodic (-1.25 and -1.375 V), Es decreases due to the higher contribution of the zinc reduction reaction. If the applied potential in made more cathodic, -1.5V, the contribution of the hydrogen evolution reaction is very important from the early stages of the experiment and, consequently, Es increases. If Figs. 3 and 5 are compared it is observed that its behaviour is reversal, as can be predicted by the comparison of equations (1) and (3).

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Fig. 3. Current efficiency of zinc vs. time as a function of the applied potential. [ZnCl2]0 = 0.055M, [HC1]0 = O.IM.

Fig. 4. Space-time yield of zinc vs. time as a function of the applied potential. [ZnCl2]o = 0.055M, [HC1]0 = O.IM.

Fig. 5. Energy consumption of zinc vs. time as a function of the applied potential. [ZnCl2]o = 0.055M, [HC1]0 = 0. IM.

In order to evaluate the iron effect on zinc electrodeposition, a synthetic solution with the same zinc and HC1 concentrations as the previous one and the same iron/zinc ratio as that present in the real sample is prepared (0.055M ZnCl2 and 0.035M FeCl2 in 0.1M HC1). Fig. 6 shows the zinc concentration evolution in the presence and absence of iron for the potential values of -1.25V and -1.5V. The solid line presents the pH profile for the most cathodic potential value of -1.5V in the presence of iron. For the potential value of -1.25V, the rate of Zn depletion is practically the same in the presence and absence of iron. However, for the potential value of -1.5 V, the zinc reduction rate becomes negligible in the presence of iron due to the great local pH increase shown in Fig. 6, that causes zinc redissolution [7], This fact is related to the enhancement of the HER process in the presence of iron, which produces a greater H+

consumption at -1.5V [13].

Fig. 6. Evolution of zinc concentration vs. time as a function of the applied potential. [ZnCl2]o = 0.055M, [FeCl2]o = 0.035M, [HC1]0 = 0.1M.

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Iron concentration was also measured and resulted invariable for all the potential values under study. This fact may be related to the high acidity of the bulk solution and to the presence of zinc that could inhibit iron deposition [6, 13]. It is noteworthy that although iron deposition does not take place, its presence in solution greatly affects the zinc current efficiency. The decrease in current efficiency observed in the presence of iron is associated with the reverse Fe+2/Fe+3 redox system taking place at the two electrodes consuming high amounts of energy [14], and to the enhancement of the HER process, which also competes for zinc electrodeposition.

Conclusions

An electrochemical reactor was developed to recover zinc from the spent pickling solutions coming from the hot dip galvanizing industry. These solutions mainly contain ZnCh and FeCh in aqueous HCl media. The different figures of merit and the product of the mass transfer coefficient and the specific surface area were calculated in order to select the best electrochemical reactor operating conditions.

The potential values to be applied in the electrolysis experiments were chosen from the voltammetric experiments and ranged between -IV and -1.75V. In the absence of iron in solution, as the electrode potential was shifted towards more negative values, the space-time yield of zinc and its fractional conversion increased because of the increase in the electrode roughness and the hydrogen turbulence-promoting action. Simultaneously, the specific energy consumption decreased initially due to the increase in the zinc conversion rate but increased for the most cathodic potential value due to HER.

Even though iron deposition does not take place for any experimental condition under study, its presence in synthetic Zn/Fe solutions led to a decrease in current efficiency associated with the reverse redox Fe2+/Fe3+ system and to the enhancement of the HER, which also induced increments in the local pH and the subsequent zinc redissolution for the most cathodic potential values.

Acknowledgements

Authors want to express their gratitude to the Universidad Politecnica de Valencia for the economical support in the project reference PAID-06-08, and to the Generalität Valenciana for the financing of the project reference GV/2010/029

References [1] U. Kerney, "Treatment of spent pickling acids from hot dip galvanizing," Resources

Conservation and Recycling, 10(1994), 145-151. [2] C. Stocks, J. Wood, and S. Guy, "Minimization and recycling of spent acid wastes from

galvanizing plants," Resources Conservation and Recycling, 44 (2005), 153-166. [3] G. Csicsovszki, T. Kékesi, and T.I. Török, "Selective recovery of Zn and Fe from spent

pickling solutions by the combination of anion exchange and membrane electrowinning techniques," Hydrometallurgy, 77 (2005), 19-28.

[4] L. Muresan et al., "Effects of additives on zinc electrowinning from industrial waste products," Hydrometallurgy, 40 (1996), 335-342.

[5] A. Recéndiz, I. Gonzalez, and J.L. Nava, "Current efficiency studies of the zinc electrowinning process on aluminum rotating cylinder electrode (RCE) in sulphuric acid medium: Influence of different additives," Electrochimica Ada, 52 (2007), 6880-6887.

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[6] M. Garcia-Gabaldôn et al., "Electrochemical study of a simulated spent pickling solution," International Journal of Electrochemical Science, 6 (2011 ), 506-519.

[7] J. Dobryszycki, and S. Biallozor, "On some organic inhibitors of zinc corrosion in alkaline media," Corrosion Science, 43 (2001), 1309-1319.

[8] F. Simescu, and H. Idrissi, "Corrosion behaviour in alkaline medium of zinc phosphate coated steel obtained by cathodic electrochemical treatment," Corrosion Science, 51 (2009), 833-840.

[9] S.C. Das, P. Singh, and G.T. Hefter, "The effects of 4-ethylpyridine and 2-cyanopyridine on zinc electrowinning from acidic sulfate solutions," Journal of Applied Electrochemistry, 27 (1997), 738-744.

[10] M. Garcia-Gabaldôn et al., "Electrochemical recovery of tin from the activating solutions of the electroless plating of polymers: Galvanostatic operation," Separation and Purification Technology, 51 (2006), 143-149.

[11] T. Boiadjieva et al., "Electrochemical studies on Zn deposition and dissolution in sulphate electrolyte," Journal of Solid State Electrochemistry, 13 (2009), 671-677.

[12] M.F. Dahab, D.L.Montag, and J.M. Parr, "Pollution prevention and waste minimization at a galvanizing and electroplating facility," Water Science and Technology, 30 (1994), 243-250.

[13] P. Diaz-Arista et al., "ZnFe anomalous electrodeposition: stationaries and local pH measurements," Electrochimica Acta, 47 (2002), 4091-4100.

[14] A.E. Saba, and A.E. Elsherief, "Continuous electrowinning of zinc," Hydrometallurgy, 54 (2000), 91-106.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

LABORATORY TESTING RESULTS OF KINETICS AND PROCESSING

TECHNOLOGY OF THE POLYMETALLIC SULPHIDE CONCENTRATE BLAGOJEV KAMEN - SERBIA

Milorad CIRKOVIC1, Zeljko KAMBEROVIC2, Vlastimir TRUJIC1,

'Mining and Metallurgy Institute Bor, Zeleni bulevar 35,Bor, Serbia, e-mail: milorad.cirkovic@irmbor.co.rs

2Faculty of Technology and Metallurgy, Belgrade, Serbia

Abstract This work presents the laboratory testing results of kinetics the oxidation process and sample processing of the sulphide polymetallic concentrate Blagojev Kamen. The aim of investigation is recovery of these types of raw material, present in large quantities in the peripheral parts of already used primary mineral deposits of copper, because of their high economic potential due to the content of a large number of metals and especially precious metals. Characterization of this raw material is based on the chemical analyses, XRD results, DTA analysis, etc. For these investigations, the sulphide concentrate with the following content was used in %: Cu - 2.3; Fe -19.8; S -27.19; Zn - 9.13; As - 0.167; Pb - 15.63; Si02 - 17.93; CaO - 0.97; A1203 - 1.43; Ag - 480 g/t; Au - 659 g/t. Kinetic investigations of oxidation processes were carried out under the isothermal conditions within the temperature range of 400 to 625 °C. The Sharp's model was used for determination the kinetics parameters, and determined values of activation energy are 67 kJ/mole for the first period, and 47 kj/mole for the second period. Pyrometallurgical treatment of this type of polymetallic concentrate, in the laboratory conditions, was carried out using the oxidative roasting and, then the reduction smelting was done in the Taman's furnace. Gold from 90.5 to 97.95% and silver from 77.28 to 93.37% are moved into the raw lead (smelting product). Gold from 1.1 to 3.92% and silver from 4.35 to 8.42% are moved into the polymetallic copper matte. Gold from 0.58 to 1.6% and silver from 2.45 to 6.82% are moved into the slag.

Keywords: polymetalic concentrate, oxidation, kinetics, Sharp

Introduction

The aim of investigation is the recovery of these types of raw material, present in large quantities in the peripheral parts of already used primary mineral deposits of copper, because of their high economic potential due to the content of a large number of metals and especially precious metals: Cu,Fe,S,Pb,Zn,Au,Ag. Goldbearing quartz Blagojev Kamen was subjected to the flotation process. Product of flotation - concentrate with content of 659 g/t Au and 480 g/t Ag. Remaining constituents in the concentrate are: As, Cd, Se, Mn, Sn, Ni, Bi, Ti, Mo, V, Co. The knowledge of mechanism and kinetics of oxidation processes of metallic sulphides is the subject of great significance in practical applications [1-3], Sulphides from the system Fe-Pb-Zn-S under oxidation processes were investigated by many researches [4-6], due to importance of those sulphides in metallurgy. However, the documentation is still incomplete and some discrepancy may be noticed by comparing different references. The starting characterisation of the investigated sample was done based on the chemical analysis, XRD results, DTA analysis as well as the constructed Phase Stability Diagrams (PSD) for system. Kinetic investigations of

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oxidation process were carried out under isothermal conditions (temperature range 400-625°C). The obtained degrees of desulphurization were used in process calculation according to Sharp's model and kinetic parameters for oxidation process were determined, including activation energies and rate constants of characteristic reactions. This polymetallic goldbearing concentrate Blagojev Kamen, with content in (%): Cu - 2.3; Fe -19.8; S - 27.19; Zn - 9.13; As - 0.167; Pb - 15.63; Si02 - 17.93; CaO - 0.97; A1203 - 11.43; Ag - 480 g/t; Au - 659 g/t was undergone to the pyrometallurgical treatment (oxidation roasting -reduction smelting in a graphite tigel in the laboratory Taman's furnace).

Experimental kinetic investigations

Thermal behaviour and mass changes analysis of starting sample, during oxidation in non-isothermal conditions were performed using the results of DTA analysis. For the purpose of the thermal analysis Derivatograph-Q (MOM, Hungary) was used with the following conditions: sensitivity DTA - lmV and DTG - 1mV, TG - 100mg, heating rate of 10°min"', sample mass lOOmg, and Trnax=1273K. A ceramic crucible was used. All experiments were carried out in the air atmosphere. X-Ray analysis was performed in order to characterize the starting investigated sample. The phase composition was determined on Siemens X-ray equipment with Cu anticathode and Ni filter at voltage 40 kV and current 18 mA. Isothermal investigations were done using an electric resistance furnace with thermostatic control. A measured volume of air was introduced into the reaction area, while gaseous product of reaction (mostly SO2), passed from furnace tube to the absorption tubes, filled with aqueous solution of hydrogen peroxide, producing sulphuric acid. The produced sulphuric acid was reacted with measured standard solution of sodium hydroxide in the presence of indicator for the purpose of calculation the sulphur content and, hence degree of desulphurization during oxidation roasting was calculated with % of expected error as ±0.1%.

Results Starting Sample Characterization

The Fe-Pb-Zn-S-0 system was interest during the investigation of polymetallic concentrate oxidation in the air. Therefore, characteristic phase stability diagrams was constructed based on starting thermodynamic data for reactions occurring in the system of interest at temperatures of 573 and 1273K, as shown in Figures 1 and 2. Phase stability diagrams were constructed for the systems Zn-S-O, Fe-S-0 and Pb-S-O separately (respectively).

Figure 1. Phase stability diagrams for the system Zn-S-O: a) 30(rC and b) 1000°C

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Figure 2. Phase stability diagrams for the system Fe-S-O: a) 30CTC and b) WOCrC

The obtained results show equilibriums between phases, with indicated final product of the oxidation process for defined concentrations of S02 and 02 in investigated systems. The results of thermal analysis of investigated sample the polymetallic concentrate sample, obtained using DTA method, are presented in Figure 3, while the results of X-ray diffraction are

presented in Figure 4.

Fig. 3. DTA analysis of the Fe-Pb-Zn-S pofymelalic concentrate sample oxidation (atmosphere - air, heating rate - I If min')

The first small endothermic peak at DTA curve, presented in Figure 3, occurs at temperature near 373K that represents a removal of humidity from sample. The following exothermal peaks are for

oxidation the sulphide samples until the corresponding sulphates are formed (FeS04, ZnSC>4 and PbSC>4). Final endothermic peak is for the sulphate dissociation to corresponding oxides and so2.

Fig. 4. X-ray ing diffractograms of the investigated startsample (Q - quartz, Ga- galena, Py - pyrite, Hkp - chalcopyrite, Sp-sphalerite According to the figure 4., dominating in the sample are well crystallised minerals from the sphalerite group mostly galena and pyrite (with sphalerite peaks partially covered with the pyrite), small amount of chalcopyrite is present as well. There is considerable amount of quarts present as well.

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Kinetics of Investigated Oxidation Process System

Degree of desulphurisation was used for determination the kinetic parameters of oxidation using the Sharp's isothermal model [7], that was later. Degree of desulphurization presents the ratio of reacted sulphur during oxidation process, compared to starting concentration of sulphur in a sample. In the course of experiments, samples were oxidized and amounts of emitted SO2 were successively registered. These results are presented in Figure 5.

Figure 5. Amount of sulfur reacted with oxygen during oxidative roasting of investigated sample (D - degree of desulfurization. r - time)

Linearization of experimental results, presented in Figure 6, was tested using nine different kinetic equations, proposed by Sharp [7], The criteria for accepting an equation as the best for linearization of the experimental data was the Sharp method of reduced half time of reaction. Using

[-ln{l-a)Y2 =k-t equation A2: for desulphurization, where: a- degree

of reaction, x-time, k-rate constant. Best linearization of experimental data was done, Figure 6. According to the Sharp's theory, equation A2 describes randomly nucleation of a new stage (Avrami I equation) [7], This stage could be concerned as a dominant in controlling the reaction rate in the first period of time. It is obvious that change in reaction mechanism occurs after first 6-12 minutes of the process what is indicated with infraction of liberalized isotherms in Figure 6. From the slopes of liberalized isotherms, the rate constants were determined and characteristic Arrhenius diagram was constructed, Figure 7.

Figure 6. Linearization of experimental data Figure 7. Arrhenius diagram for the process of points using kinetic equation A 2 for the first oxidative roasting constructed according to the stage of the desulphurization process slope of sulphur removal during the process

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According to the Arrhenius diagrams, the activation energy of process, under isothermal conditions was calculated. Activation energy is 67 kj/mol for the oxidation process before infraction of liberalized isotherms and 47 kj/mol after it.

Laboratory Pyrometallurgical Investigations

Today, in the world, for gold extraction from sulphide ores and concentrates, the pyrometallurgical methods are used more and more, and the technologies as cianization and amalgamation [8,10] are left. The basic reason for this is in higher recovery of gold and dificulties in use and living and environment protection in work with cyanides.

In the process of pyrometallurgical extraction of precious metals there is always the basic phase - collector, where precious metals are collected. Selection of collector phase depends on content the starting raw material, further metallurgical treatment and other factors. In metallurgical treatment of suphide gold bearing concentrates as collector, the most usable are lead, copper or calcine.

Lead is traditionally used as a collector for precious metals due to high dissolubility of precious metals in lead at low temperatures, rapid separation of lead by pyrometallurgical method (cupellation process) and low working temperature (up to 1000°C). This is especially if goldbearing concentrate or secondary raw materials contain significant quantity of lead. The main shortage of lead as collector is the required rigorous control of environment pollution and unsatisfactory recovery of I, Ru and Os in cupellation phase [9] .

In decision making on selection the collecting phase in certain situation, besides the above mentioned, the other facts have to be taken into consideration as well as : chemical content of concentrate, what in some cases could have a decisive influence, consumption of normative material, training of labour force and others.

Results of Experimental Investigations

Laboratory experimental investigations were carried out in the laboratories of Mining and Metallurgy Institute Bor on a concentrate sample from Blagojev Kamen with the following content (%): Cu - 2.30; Fe - 19.80; S - 27.90; Zn - 9.13; As - 0.167; Pr - 15.63; Si02 - 17.93; CaO - 0.97; A1203 - 1.43; Ag - 480 g/t; Au - 659 g/t.

Investigations were carried out by smelting the concentrate of 1 kg concentrate in a graphite tigel in the laboratory Taman's furnace. First investigation was carried out by smelting of unroasted concentrate sample. Two phases were obtained: calcine and slag. Calcine had the following content (%): Cu - 6.81; Pb - 9.00; Zn - 9.31; S - 30.80; Fe - 40.49; Au - 0.129 and Ag - 0.094. Content of calcine was a little bit different from concentrate. Slag was very acidic with the following content: SiOz - 58.0; CaO - 12.07.

This method of concentrate smelting does not give the satisfied results because the polymetallic Cu - Pb - Zn calcine is obtained with about 60% of transfered gold and silver.

The following série of four samples was carried out according to the following layout: oxidation roasting - smelting with addition of reducer and without reducer in presence of

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suitable fluxes. Three investigations were carried out with roasting material with content of 7% S, and one investigation with roasting material with 3.55% S. Smelting of roasting material was carried out with variable quantity of reducer (powdered coal). Smelting of samples of 1 kg roasting material was carried out in the laboratory furnace in a graphite tigel at 1250°C. Graphite tigel was cooled by various intensity. Hardened melt includes three layers - raw lead on the bottom of tigel, Cu-Pb-Zn calcine in the medium part and slag on the suface.

Based on the results of laboratory investigation and literature data, the following technological layout (Figure 8) was defined for treatment the sulphide polymetallic concentrate from Blagojev Kamen: oxidation roasting, charge preparation, smelting, cupellation of raw lead to dore metal and refining of dore metal. The first three phases of process were checked in the laboratory conditions (oxidation roasting of concentrate, charge preparation and smelting).

Figure 8. Block sheme of treatment the goldbearing concentrate from Blagojev Kamen

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Metal part is raw lead with content of 91.32 to 96.38% Pb and, by content, lead is very similar to raw lead, produced in the lead metallurgy in a pit furnaces. In our case, content of precious metals, especially gold, is much more higher, in the range of 0.66 to 1.38%, and silver of 0.598 to 1.11%, since in raw lead, produced in the lead metallurgy, content of those metals is much more lower, 1000 - 5000 g/t Ag, 1-5 g/t Au.

Second phase of smelting process is presented by Cu - Pb - Zn calcine that is also similar to calcine, produced in the lead metallurgy in a pit furnaces. Calcine contains: 15 - 24% Pb; 13.77 - 42% Cu; and 2 - 4% Zn. Beside this, it also contains of 0.02 to 0.14% Au and 0.22 to 0.8% Ag.

Third stage of slag smelting has the following content: SiCh - 36.45%; FeO - 26 - 31%; and CaO - 10 - 20%; and minimum quantity of useful metals: 1 to 1.55% Pb; 0.16 to 1.39% Cu; about 5% Zn; 0.0006 to 0.001 Au; and 0.0055 to 0.25% Ag.

Transfer degree of useful metals, especially precious metals into raw lead is very high: Pb from 50.23 to 55.88%; Au from 90.5 to 97.95% and Ag from 77.28 to 93.37%. Raw lead, produced in the investigation process, is very good collector for precious metals and, in a case of good smelting process control, nearly total quantity of precious metals is transfered into raw lead.

In the Cu-Pb -Zn, the following is transfered: 4.48 to 10.47% Pb; 52.35 to 74.37% Cu; 1.1 to 3.92% Au and 4.35 to 8.42% Ag. Calcine presents a reverse material where additional metal recovery could be carried out by supplemental treatment. Unsignificant quantity of metals is transfered into slag: 3.88 to 8 % Pb; about 0.58 to 1.6% Au and 2.45to 6.28% Ag. Very low is a transfer degree of gold into slag.

Conclusion

Investigations the new copperbearing raw materials, which are located on the marginal parts of the main exploited mineral deposit of copper ore, are significant because of the upcoming modernization of the Copper Smelter in RTB Bor.

Goldbearing concentrate with content in (%) of : Cu - 2.3; Fe - 19.8; S - 27.90; Zn -9.13; As - 0.167; Pb - 15.63; Si02 - 17.93; CaO- 0.97; was used for laboratory investigations of pyrometallurgical process the extraction of precious metals. Goldbearing raw material from the locality Blagojev Kamen with content in (%) of : Cu - 0.047; S - 0.99; Pb - 0.11; Zn - 0.049; Si02 - 87.14, AI2O3 - 1.98; Au - 10-15g/l; Ag - 15-25 g/1; was undergone to the flotation process in the semi-industrial plant of Mining and Metallurgy Institute Bor.

The investigation results of oxidation process in the system Fe-Pb-Zn-S are presented in this paper. In the first part of paper characterisation the starting sample, including chemical analyses, X-ray diffraction, as well as thermodynamic analysis of phase stability in the characteristic systems, are presented. The process kinetics was investigated using the Sharp method and it was determined that the oxidation process starts rather fast at the beginning of the process (first 6-12 minutes, depending on temperature) with activation energy which places the process in kinetic region. After this first period, activation energy fall to the value on border of kinetic-to-transition area, meaning that diffusion of reactants and product of reaction began to have influence on overall rate of reaction.

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Based on the laboratory investigation results and literature data, the following technological layout was defined for treatment the sulphide polymetallic concentrate from Blagojev Kamen: oxidation roasting, charge preparation, smelting, cupellation of raw lead to dore metal and refining of döre metal. The first three phases of process were checked in the laboratory conditions (oxidation roasting of concentrate, charge preparation and smelting).

The investigations of technology for goldbearing quartz from the locality of East Serbia were carried out in the laboratories of Mining and Metallurgy Institute Bor. Goldbearing quartz was subjected to the flotation process. Product of flotation- concentrate with content of 659 g/t Au and 480 g/t Ag was subjected to the pyrometallurgical treatment: oxidation roasting, reduction smelting.

Three stages were obtained in the smelting process: raw lead with content of 91.21 to 96.38 % Pb; 0.74 to 0.809 % Au and 0.49 to 0.598 % Ag. Cu - Pb - Zn calcine with content of 15.25 to 24.76 % Pb; 13.77 to 41.99 % Cu; 2 - 4 % Zn; 0.02 to 0.075 % Au and 0.036 to 0.128 % Ag. Slag with content of 0.0006 to 0.0013 % Au.

The majority of precious metals from concentrate are transformed into raw lead as well as: 90.5 to 97.95 % Au and 77.28 to 93.37 % Ag.

The proposed technology for goldbearing quartz treatment consists of the following stages: - flotation of previous crushed quartz ore, - oxidation roasting of flotation concentrate, - reduction smelting, - cupellation of raw lead even to dore metal, - refining of döre metal.

References

1. G.V.Samsonov, S.V.Drozdova, Sulfides, Metallurgy, Moscow, 1972. (in Russian) 2. D.J.Vaughon, J.R.Craig, Mineral chemistry of metal sulfides, Cambridge University Press,

Cambridge, 1978. 3. N.Strbac, D.Zivkovic, Z.Zivkovic, I.Mihajlovic, Monograph: Sulphides - thermal,

thermodynamic and kinetic analysis, Punta, Bor, 2005. (in Serbian) 4. A.W.Espelund and H. Jynge, Scandinavian Journal of Metallurgy, 6(6) (1977), 256-262. 5. V.M.Kuzminykh et.al., Izvestiya Vysshikh uchebnykh Zavedenij. Tsvetsnatya Metallurgiya,

2(1984)59-64. 6. D. Sinadinovic, R. Vracar, Z. Kamberovic, CIM Buletin, 94 (1051) (2001) 123-128. 7. H.J.Sharp,W.G.Brinduy.N.B.Harahari, J.Amer.Ceram.Soc, 49 (1966) 379. 8. G. Hilson, AJ. Monhemius, Alternatives to cyanide in the gold mining industry:

What prospects for the future?, Journal of Cleaner Production 14 (2006) 1158-1167 9. Gold Amalgamation, http:www.goldDrospecting.com 10. K. Tanida, M. Hoshino, Continuous Determination Of Mercury Air By Gold

Amalgamation And Flameless Atomic Absorption, The Rigaku Journal, Vol. 7, N°2, 1990, 35/40.

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Hidrotalcite with Gentamicine, of the Type Mg0.68Alo.32(OH)2(N03)o.32,0.1H20, Formed by Chemical Coprecipitation in Controlled Atmosphere

H. H. Rodriguez ', O. Martinez-Alvarez1.

1. Universidad Politécnica de Guanajuato, Av. Universidad Norte S/N, Juan Alonso Cortazar Gto. C.P. 38438

* E-mail of the author: ekkodnam@gmail.com, Phone: +524614414300, Fax: +524614414328.

Abstract This work consisted in the development for coprecipitation and characterization for X-ray diffraction of a hidrotalcite of the type Mgo68Alo32(OH)2CN03)o32*0.1H20, capable of liberate in vitro gentamicin. The objective of this project was to generate a new material capable of remedy infections for the bacteria Staphylococcus Aureus, when this bacterium is introduced fortuitously in the body of a patient through of implants o prosthesis. A new method at the vacuum for condensation of vapor was developed of the hidrotalcite synthesis. Gentamicin was introduced in ionic clay, for a process of ionic interchange. The sensibility to the new product was tested in Staphylococcus aureus antibiograms. The results suggest a high sensibility of this kind of bacteria to the new product.

Keywords Ceramics, biomaterials, hidrotalcite, drug delivery, implants.

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Introduction

In orthopedic and traumatology are used prosthesis and implants of alloys considered like bioacceptable. These replace or give support to the bony system when this one has undergone a severe trauma that affects its performance operation [1> 2l It is required that it implants or prosthesis has the mechanical, physical and chemical properties that do not affect the functionality of the adjacent tissue and that their performance is equivalent or superior to the part that replaces [ M l An additional requirement is that these devices are not an infectious vector '5'7'. In spite of these requirements, this has been observed that appear infection in patients, between the 1 and 2% of all the prosthesis implanted [3,91, causes conditions of risk for the patient and stages of pain, that only can be remediated for chirurgical intervention [5].

A preventive alternative to avoid the infection can be the use of controlled liberation medicine [5]. Nevertheless, if it is use combined with acrylic cements can diminish the efficiency of these exposed medicines to exothermic conditions [5]. Other cements developing are the alpha tricalcic phosphate (a-TCP) based. That has the capacity of transformed into the mineral phase of the bone [10l

The possibility exists of using hidrotalcite (HT) with medicine adhered to its surface. Which can be integrated to a biocement based calcium phosphate, which provides endothermic conditions that once fixed between the treated bone and it implants can be transformed into the mineral phase of the bone and while this happens can be left the medicine impregnated into HT, allowing its liberation free. One of the considered HT as bioacceptable is the Mgo 6sAlo,32(OH) 2(N03)o 32*0.1^0 which has been impregnated of folicacid1"1

Material and methods

The used materials was the following ones: deionized water, nonahidrated aluminum nitrate hexahidrated, hydroxide magnesium nitrate of sodium, injectable gentamicin of 80mg. Stocks of Staphylococcus control aureus (Gram+, S. aureus ATCC® 29223), Agar Mueller-Hinton, defibrinated blood of ewe 5% (v/v), and discs of filter for test of diffusion.

The hidrotalcite was prepared by coprecipitation with a controlled atmosphere [ " l Mgo 68Alo.32(OH)2(N03)o 32-0.1H20 was prepared by coprecipitation under vacuum atmosphere. A mixed solution of Mg(N03)2*6H20 (0.032 m) and A1(N03)3*9H20 (0.016 m) and it was titrated with NaOH (0.5 m) solution at room temperature. The final pH of the solution was adjusted to 9.570.2. The resultingwhite precipitate was collected by centrifugation, and washed with deionized water thoroughly. For what the HT acquires the medicine, the HT was dispersed in a solution that contained in gentamicin in excess. It was shaken intensely to 60°C by 3 days. The products were washed and they were dried to the emptiness, and they were kept in a sterilized desiccators waiting for its use.

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In order to establish the sensitivity of the Staphylococcus aureus to the new product, tests of the type antibiograms were made. Discs of filter paper were used, the HT with antibiotic was compact to pressure on the discs, of this way a structure occurred him to the ceramics and the method of diffusion by disc was simulated (Kirby-Bauer method). These discs were used later depositing itself in the plates of Agar[l3'151.

The sterilization of instruments, the preparation of the Agar Mueller-Hinton, monitoring of pH and the humidity, as well as the standard of turbidity for preparation of inoculate were through standard methods[13"151 already proven.

For the preparation of the inoculate was used the direct method of suspension of colonies. Inoculate was prepared directly making a broth of isolated colonies of a plate of agar of 18 to 24 hours (agar-blood). The suspension was fit up to 0,5 McFarland of turbidity. The inoculation of the plates followed the following process: fifteen minutes after fitting the turbidity of the suspension of the inoculate, with a hyssop of cotton was absorbed somewhat of the same one suspension. Hyssop was rotated several times and pressed firmly against the internal wall of the tube on the liquid level to eliminate the excess of the inoculate. The surface of the plates of agar Mueller-Hinton was inoculated whit the hyssop on the entire surface. This procedure was repeated lining two or more times, turning the plate approximately 60°C every time to assure a constant distribution of the inoculate and in the end the edges of agar inoculated. The plates left minutes opened to allow that an excess of humidity was absorbed before applying the disc with the HT.

For the application of discs of Ht with antibiotic, the discs with HT and antibiotic were placed on the surface of agar with the face with HT and antibiotic towards agar having been avoided their relocation and neighborhoods among them smaller to 30 mm. After applied to the discs the plates they were invested and put in an incubator 37°C. This procedure was used to inoculate a plate of agar with a disc with HT without antibiotic like control. For the interpretation of the halos of inhibition, after 18 hours of exposition of the plates to agar to the HT with gentamicin, the halo formed was measured in millimeters using a digital vernier. The rank of reference of susceptibility according to the disc method Kirby Bauer was in this project of: Sensible: >18 mm, intermedium: 15 - 17mm, resistant: <14 mm.

By scanning electron microscope (SEM) the HT was observed. By X-ray diffraction (XRD) and Energy-dispersive X-ray spectroscopy (EDX) the nature of formed compounds was determined. The efficiency of the antibiotic was determinate measuring the halo produced by the HT with antibiotic. As control were used a HT without antibiotic.

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Results

The preparation of the HT, following the proposed method, gave by result a very fine dust of white color as it is observed in figure 1. The size of particle varied of 0,3 to 2 microns, appearing in general like cluster and not like individual particles.

The pattern of the Ht resulted for analysis by XRD was of the kind Mgo68Alo,32(OH)2(N03)o32*0.1H20 of figure 2. After to be immersed in the solution mother with the antibiotic, it continued conserving this color whereas his pattern of XRD changed when appearing a tip located in position 5 of the axis 20, which indicates the integration of the gentamicin between the interlayer is of the hidrotalcite. This change is in the pattern of figure 2.

Agar obtained and the seeded cultures they did not present problems, reason why it presented a standard quality in the samples.

The test of exposition of the HT without gentamicin to the Staphylococcus aureus is showed in figure 3. It is possible to be observed that the HT did not interfere in the growth of the culture, reason why it is possible to be considered that this material does not have biocides properties that obstruct in the efficiency of the antibiotic.

The HT with gentamicin were used altogether in 56 discs, and exposed individually to cultures of Staphylococcus aureus obtaining the results that are described in table 1.

The histogram formed with the test resulted of sensitivity to the antibiotic with HT [figure 4], shows that 94,6% of discs with HT and gentamicin had a biocide effect on the culture of Staphylococcus aureus causing of inhibition halos of growth of the culture with a diameter average of 28,6 mm.

Discussion

When fixing mechanically by pressure the HT with gentamicin to discs for the test of sensitivity to the antibiotic, was observed that the dust was adhered to the disc depending on the care whereupon was handled since showed certain plasticity when it was compacted.

The obtained thicknesses were not measured for lacking the appropriate equipment to do it, a constant thickness was supposed using constant units of mass, pressure and area.

The significance test was so that is not required to test of given hypothesis of according to the standard deviation. The slant of the curve can be interpreted like an effect caused by the activity of the antibiotic since it was still active, as shown in Figure 4.

5.4% of discs with HT and gentamicin showed an intermediate sensitivity. These discs were the last ones in being placed and with the most recent HT in agar by a new assistant. Possibly several factors took part in this difference of sensitivity, among them the factor of pressure differential when being placed the discs on agar between first 53 discs and these last 3 find repercussions in the effect of the antibiotic.

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The effect of the antibiotic is respectively observed in the image of the figure 5a and 5b. In the results with intermediate sensitivity (Figure 5b), perhaps, the longevity of the culture front a smaller time of exhibition could have taken part since in these last 3 discs a time of exhibition of 16 hours was used. The figure 5a shows halos of a greater diameter and some new growth of the bacterium was not observed in them.

Conclusions.

The procedure to obtain the HT of the type Mgo68Al0>32(OH)2(N03)o32*0.1H20 and the form to integrate the antibiotic to the ionic clay were the correct ones to obtain an ionic ceramics with the capacity to catch and to release antibiotic that inhibits the growth of bacteria of the Staphylococcus aureus in agar plates. Reason why the use of this type of material can be used in medicine.

Acknowledgements

The author thankful widely, to the CONACYT, by the financing of his post Doctorate, and to the PROMEP for the financial support.

References 1. Jorgensen J. H. Laboratory issues in the detection and reporting of antibacterial resistance. Infectious Disease Clinics of North America ( 1997); 11: 785-802. 2. Jones RN, Biendenbach Dj, Marshall SA, Pfaller MA and Doern GV Evolution of Vitek System to Accurately Test the Susceptibility of Pseudomonas Aeruginosa clinical isolates against Cefepime. Diagn Microbiol Infect Dis, (1998); 32: 107-110. 3. Ferraro M. J, Jorgensen J. H: Instrument-based antibacterial susceptibility testing. In Murray PR, Baron Ej, Pfaller MA et al (eda): Manual of Clinical Microbiology, Washington, DC, American Society for Microbiology, (1995), pp 1379-1384. 4. B. M. Sanchez, E. Delgado, A. Perera and C. Evora, In vitro-in vivo characterization of gentamicine bone implants, J. Control Release, vol 83, (2002), 353-364 5. Alex C. McLaren, Alternative Materials to Acrylic Bone Cement for Delivery of Depot Antibiotics in Orthopaedic Infections, Clinical Orthopaedics and Related Research ,427,(2004),101-106 6. Li XD, Hu YY: The treatment of osteomyelitis with gentamicin reconstituted bone xenograft-composite. J Bone Joint Surg, 83B, (2001), 1063-1068. 7. Nazarov AD, Firsov AA, Navashin PS, Fomina IP, Rudenko TG: Pharmacokinetic study of implantable gentamycin preparations. I. Antibiotic pharmacokinetics in the implantation area and an evaluation of the prolonged effect of the preparations. Antibiot Khimioter 33, (1998), 605-612. 8. Ostermann PA, Seligson D, Henry SL: Local antibiotic therapy for severe open fractures. A review of 1085 consecutive cases. J Bone Joint Surg 77, (1995), 93-97. 9. Gitelis S., GT Brebach G.T.; The Treatment of Chronic Osteomyelitis With A Biodegradable Antibiotic Impregnated Implant; Journal of Orthopaedic Surgery; (2002); 10: 53-60. 10. Schmidmaier G, Lücke M, Wildemann B, Haas NP, Raschke M., Prophylaxis and

treatment of implant-related infections by antibiotic-coated implants: a review. Injury,

37,(2006), 105-112

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l l . J . Choya, J. Junga, J. Oha, M. Parka, J. Jeong, b. Kangb, O. Hanb, Layered double hydroxide as an efficient drugreservoir for folate derivatives. Biomaterials 25 (2004) ,3059-3064 12. Khairoun I, Boltong MG, Driessens FCM, Planell JA; Effect of Calcium Carbonate on The Compliance of Apatitic Calcium Phosphate Bone Cement; Biomaterial; 18, (1997), 1535-1539. 13. N. Bejoy, Hidrotalcites, the clay that cures, Resonance, (2001), 57-61. 14. National Committee for Clinical Laboratory Standards, NCCLS, (1998). Performance Standards for Antimicrobial Susceptibility Testing, 8th International Supplement M 100 S8, Wayne, PA: NCCLS. 15. Rittenhouse SF; Miller LA; Utrup LJ; Poupard JA. Evaluation of 500 gram negative isolates to determine the number of major susceptibility interpretation discrepancies between the Vitek and Microscan Walkway for 9 antimicrobial agents. Diagn Microbiol Infect Dis, (1996); 26 [1]: 1-6. 16. Biendenbach DI, Marshall SA and Jones RN. Exactitud de los resultados de la Susceptibilidad a los Antimicrobianos para Pseudomonas aeruginosa evaluada por el sistema Microscan WalkAway, Diagn Microbiol Infect Dis (1999); 33: 305-307.

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Figures and tables

Figure 1. Micrograph of the HT obtained by coprecipitation.

Figure 2. XRD Difractograms of the HT and the HT with gentamicin respectively.

Figure 3. Antibiogram of the disc with HT without antibiotic.

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Table 1. Statistical of the results of antibiograms, mark [I] indicates an intermediate sensitivity whereas mark [S] indicates a reliable sensitivity to the antibiotic.

Figure 4. Histogram of frequency of the sensitivity of the Staphylococcus Aureus to the HT with gentamicin.

Figure 5. Images on the same scale of antibiograms, sensible [a] and intermediate sensitivity [b] obtained when exposing discs with HT and antibiotic to the inoculated

Agar.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF THIODIGLYCOLAMIDE ADDITION TO Di-n-HEXYL SULFIDE ON THE Pd(II) EXTRACTION RATE

Hirokazu Narita1, Mikiya Tanaka1 and Shinji Ueno2

'Research Institute for Environmental Management Technology, National Institute of Advanced Industrial Science and Technology (AIST)

16-1 Onogawa, Tsukuba, Ibaraki 305-8569, Japan 2Catalysts Development Center, N. E. CHEMCAT, Japan

Keywords: Palladium, Solvent Extraction, Thiodiglycolamide, Catalytic Effect

Abstract

Di-n-hexyl Sulfide (DHS) has been widely used as the practical Pd(II) extractant. Although its selectivity of Pd(II) over Pt(IV) is high enough in an acidic chloride solution, the extraction is very slow. In order to accelerate the Pd(ll) extraction while retaining its high selectivity and operability, one method is to add an acceleration reagent to the organic phase. We have found that thiodiglycolamide (TDGA) can immediately extract Pd(II) in hydrochloric acid solutions. Therefore, in this study, we investigated the effect of the TDGA addition on the extraction rate of Pd(II) from hydrochloric acid solutions. In the 0.1 M DHS (diluents: 80 vol% w-dodecane and 20 vol% 2-ethylhexanol)—5 g/L of Pd(II) in 1 M HCl system, it takes over 10 hours to attain the extraction equilibrium. In contrast, the Pd(II) extraction is accelerated by adding a small amount of AyV'-dimemyl-A'.AMi-w-octyl-thiodiglycolamide (MOTDGA); the extraction percentage of Pd(II) with 0.001 M MOTDGA—0.099 M DHS reached almost 100% in a short extraction time (-60 min). This suggests that MOTDGA has a catalytic effect on accelerating the Pd(II) extraction with DHS.

Introduction

Among both the metal mining and recycling processes of the platinum group metals (PGM), the separation and purification of PGM are mainly performed by solvent extraction. In particular, the separation flow at the INCO refinery is one of the most popular PGM separation processes, in which PGM is first dissolved using concentrated chloride media, and then individually separated by distillation, solvent extraction, etc.[l, 2] In this process, the palladium extraction is performed using dialkyl sulfide (DAS) (e.g., di-n-hexyl sulfide (DHS) and di-n-octyl sulfide (DOS)).

Although DAS can selectively extract Pd(II) over Pt(IV) from the HCl solution, its Pd(II) extraction is very slow. According to the HSAB rule [3], the Pd(II) ion that belongs to soft acids has a high affinity for a sulfide of DAS classified as soft donors. However, the sulfide is readily oxidized to a sulfoxide, which leads to a decrease in the Pd(II) extraction [4, 5],

We have studied N, TV-disubstituted amide compounds as new extradants for PGM and gold, and found a few efficient extradants [4-10]. Among them, thiodiglycolamide (TDGA) compounds allow the rapid extraction of Pd(II) and a high oxidation resistance compared to DHS [4-6], Therefore, TDGA compounds are promising as practical extradants.

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Presently, DHS is commercially available, while TDGA has not yet been widely circulated. Hence, the production cost of TDGA would be higher than that of the conventional extradant, DHS. Therefore, in the present study, we carried out the Pd(II) extraction from an HCl solution using mixed extradants in which a small amount of JVyV-dimethyl-JvyV-di-n-octyl-thiodiglycolamide (MOTDGA) was added to DHS (Fig. 1) in order to investigate the effect of the MOTDGA concentration on the Pd(II) extraction rate. Back extraction of Pd(Il) with an ammonia solution was also studied.

Figure 1. Structure of extradants (A) MOTDGA, (B) DHS

Experimental

Reagents

The stock solution was prepared by dissolving the Pd and Pt chlorides, (PdCb, H2PtCl6: Soekawa Chemical Co., Ltd.) into HCl solutions. DHS was purchased from Wako Pure Chemical Industries, Ltd. MOTDGA was synthesized using a procedure reported in reference 6. All the other chemicals used in this study were of reagent grade.

Extraction and analytical procedures

DHS and/or MOTDGA diluted by 80 vol% «-dodecane—20 vol% 2-ethylhexanol was pre-equilibrated with the same volume of an HCl solution in the absence of the metal ions. One milliliter of the pre-equilibrated organic phase and the same volume of the HCl solution containing 5 g/L each of the mixed metal ions (Pd(II), Pt(IV)) were vertically shaken in a 10-mL glass tube at an amplitude of 100 mm and a frequency of 200 spm for the required time, then centrifiiged. The back extraction experiments were performed as follows: a portion of the organic phase was transferred to another glass tube, and a same volume of a 28% ammonia solution was added; the mixture was shaken for 60 min and the metal ions in the organic phase were back-extracted into the aqueous phase. The volume ratio of the organic phase/aqueous phase {01 A) was 1 during all the experiments. All the extractions were carried out at room temperature (23 ±

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2 °C). The concentrations of the metal ions in the aqueous phase were measured by ICP-AES (Horiba ULTIMA2); those in the organic phase were calculated on the basis of the mass balance of the metal ions before and after the extraction.

Results and Discussion

Separation of Pd(II) over Pt(IV)

In all the systems (DHS, MOTDGA and MOTDGA—DHS mixed solvents), the extraction percentage (E%) of Pd(II) was nearly 100 when the extraction equilibrium was attained in each system. On the other hand, Pt(IV) in each extraction system was hardly extracted.

Dependence of Pddll extraction rate on the addition of MOTDGA to DHS

Figure 2 shows the extraction of Pd(II) from 1 M HC1 as a function of the extraction time using DHS, MOTDGA and MOTDGA—DHS mixed solvents. The total concentration of the extractants ([MOTDGA] + [DHS]) was 0.1 M. In the DHS system, the E% of Pd(II) was about 75% after the 240-min extraction. It takes over 10 hours to reach almost a 100% extraction. In contrast, 0.1 M MOTDGA (DHS absent) extracts almost all the Pd(II) ions within 10 min. In the MOTDGA—DHS systems, the Pd(II) extraction is much faster than that in the DHS system. The Pd(II) extraction pattern using the 0.005 M MOTDGA—0.095 M DHS and 0.01 M MOTDGA—0.09 M DHS mixed solvents is almost the same as that using 0.1 M MOTDGA (DHS absent). The E% of Pd(II) with 0.001 M MOTDGA—0.099 M DHS exceeded 99% after the 60-min shaking.

Figure 2. Extraction of Pd(II) from 1 M HC1 as a function of extraction time using DHS, MOTDGA and DHS—MOTDGA mixed solvents. ([MOTDGA] + [DHS]): 0.1 M. [Pd]: 5 g/L.

Since the PdCU2" ion dominantly exists in the 1 M HC1 solution [11] and the DHS and MOTDGA are neutral extractants, the Pd(II) extraction equilibria with DHS and MOTDGA are likely to be expressed as:

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PdCU2- (.,) + 2DHS(org) «-> [PdCl2(DHS)2](or8) + 2Cr(aq)

(1)

PdCl42"(aq) + 2MOTDGA ( o r g )^ [PdCl2(MOTDGA)2](org) + 2Cl"(aq) (2)

where (aq) and (org) denote the species in the aqueous and organic phases, respectively. The ratio of Pd(II) to extractant is 1:2 in both systems. In this study, we used 5 g/L (-0.05 M) of Pd(ll) and 0.1 M of the extradants (sum of [MOTDGA] and [DHS]), indicating that free extradants, which are not involved in the extracted complexes in the organic phase, are hardly present when the E% of Pd(II) is nearly 100%.

At the 120-min extraction, the E% of Pd(II) with 0.1 M DHS is about 70%, which has not equilibrated yet; on the other hand, that with 0.001 M MOTDGA—0.099 M DHS is almost 100%. This means that DHS and MOTDGA in the 0.001 M MOTDGA—0.099 M DHS system should extract -0.035 M Pd(II) and -0.015 M Pd(II), respectively, assuming that the DHS and MOTDGA separately extract Pd(II). However, 0.03 M MOTDGA is required in order to completely extract 0.015 M Pd(II) due to the formation of the 1:2 Pd(II) to MOTDGA complex. Since the concentration of MOTDGA is 0.001 M, the MOTDGA molecule is very insufficient for the Pd(II) extraction. This suggests that the MOTDGA molecules catalyze the acceleration of the Pd(II) extraction with DHS.

Back extraction of Pd(Il)

The back extraction of Pd(II) in the organic phase was carried out using a 28% ammonia solution. In the 0.1 M MOTDGA system (DHS absent), the Pd(II) is hardly back-extracted. In contrast, the Pd(II) can be back-extracted in the other systems: -70% in the 0.01 M MOTDGA—0.09 M DHS system, -80% in the 0.005 M MOTDGA—0.095 M DHS, 0.01 M MOTDGA—0.09 M DHS and 0.1 M DHS systems (MOTDGA absent).

Conclusions

We investigated the extraction properties of Pd(II) and Pt(IV) in an HC1 solution when a small amount of MOTDGA was added to DHS in 80 vol% «-dodecane—20 vol% 2-ethylhexanol. The addition of 0.001 M MOTDGA to 0.099 M DHS drastically increased the Pd(II) extraction rate. The MOTDGA molecule worked as a catalyst with respect to accelerating the Pd(II) extraction. The back extraction of Pd(II) in the organic phase can be performed using a 28% ammonia solution in the DHS and MOTDGA—DHS systems.

Acknowledgment

A part of this work was financially supported by Industrial Technology Research Grant Program from New Energy and Industrial Technology Development Organization (NEDO) of Japan.

References

1. J. E. Barnes and J. D. Edwards, "Solvent Extraction at INCO's Acton Precious Metal Refinery," Chemistry and Industry, (1982), 151-155.

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2. M. Cox. "Solvent Extraction in Hydrometallurgy," Principles and Practices of Solvent Extraction, ed. J. Rydberg, C. Musikas and G.R. Choppin (New York, NY: Marcell Dekker, Inc., 1992), 357-412.

3. R. G. Pearson, "Hard and Soft Acids and Bases," Journal of American Chemical Society, 85 (1963), 3533-3539.

4. H. Narita, M. Tanaka and K. Morisaku, "Durability of two extradants for Pd(II) separation, thiodiglycolamide and di-n-hexyl sulfide: Against a mixed solution of HNCh and HC1," Sohn International Symposium Advanced Processing of Metals and Materiah, Vol. 4, (2006), 589-594.

5. H. Narita, M. Tanaka and K. Morisaku, "Palladium Extraction with N,N,N',N'-Tetra-n-OctyI-Thiodiglycolamide," Minerals Engineering, 21, (2008), 483-488.

6. H. Narita et al., " Rapid Separation of Palladium(Il) from Platinum(IV) in Hydrochloric Acid Solution with Thiodiglycolamide," Chemistry Letters, 33, (2004), 1144-1145.

7. H. Narita et al., "Structural Effect of Monoamide Compounds on the Extraction of Gold," Solvent Extraction Research and Development, Japan, 12, (2005), 123-130.

8. H. Narita et al., " Extraction of Gold(IIl) in Hydrochloric Acid Solution Using Monoamide Compounds," Hydrometallurgy, 81, (2006), 153-158.

9. H. Narita et al., " Extraction of Platinum(IV) in Hydrochloric Acid Solution Using Diglycolamide and Thiodiglycolamide," Solvent Extraction Research and Development, Japan, 13,(2006), 101-106.

10. H. Narita, K. Morisaku and M. Tanaka, "The First Effective Extradant for Trivalent Rhodium in Hydrochloric Acid Solution", Chemical Communications, (2008), 5921-5923.

U . S . A. Wood, B. W. Mountain and P. Pan, "The Aqueous Geochemistry of Platinum, Palladium and Gold — Recent Experimental Constraints and a Réévaluation of Theoretical Predidions," Can Mineral, 30, (1992), 955-982.

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Synthesis and characterization of metallic oxides

Eduardo Brocchi1, Marina Doneda1, Rogério Navarro1, Ana Cristina Wimmer1, Rodrigo Souza1, José Campos2

'PUC-Rio (Pontifical Catholic University of Rio de Janeiro); Rua Marques de Säo Vicente, 225 - Gâvea; Rio de Janeiro, RJ, 22451-900, Brazil

2UERJ (Rio de Janeiro State University); Rua Fonseca Teles, 121 - Sào Cristöväo; Rio de Janeiro, RJ, 20940-200, Brazil

Keywords: synthesis, characterization, metallic oxides, nitrates dissociation

Abstract

Due to a recent interest in nanostructured materials this work proposes an alternative synthesis method as compared to the already established aqueous phase precipitations and gaseous compounds reactions. The method is based on the pyrolisis of Fe, Cu, Ni, Zn and Al nitrates, including a theoretical thermodynamic analysis and the corresponding experimental characterizations of the products, by XRD and SEM. Thermodynamics predicts that the decomposition reactions are possible in the temperature range between 150°C and 600°C. XRD patterns showed the formation of Fe2C>3, CuO, NiO, ZnO and AI2O3. SEM indicates the occurrence of a particle cluster structure for Fe, Cu and Ni oxides while the two others are formed either as faceted crystal (ZnO) or in the pancake shape (AI2O3). Further analysis by the Rietveld method proved that this methodology can be used to produce oxides containing nanosized particles.

Introduction

The nanostructured materials still attracts the attention of the scientific community once numerous studies have been carried out in order to improve the knowledge about their behavior and particular properties. The efforts to study this kind of materials are related to the technological potential associated with them as many segments of the industry may benefit from their outstanding novel properties (e.g. electronics, energy, aerospace). This actual context shows the magnitude of these researches, the role of the nanostructured materials in the future as well as the necessity of development alternative synthesis methods as compared to the already established aqueous phase precipitations'1' and gas-gas reaction routes'2'. Several studies are being performed to obtain nanotubes (e.g. carbon, titania)'3'4' and other nanostructured compounds (e.g. oxides)'5'. These oxides, in addition to their eventual directly applications, can also be used as precursors to obtain materials of a different chemical nature (e.g. alloys, composites)'6'. Therefore, the main objective of this work is the synthesis of metallic oxides of Fe, Cu, Ni, Zn, and Al, based on the thermal decomposition (pyrolisis) of their respective nitrates, followed by their characterizations through x-ray diffraction, as well as scanning electron microscopy. In order to better identify the operational conditions for the oxide synthesis, a thermodynamic analysis was applied in the study of the nitrate dissociations viability.

179

Thcrmodynamic considerations

The thermodynamic viability of the nitrates thermal decomposition can in a first approach be appreciated by plotting the Gibbs free energy of reaction as a function of temperature (Figure 1).

Figure 1. AG° (kcal/mol) vs. Temperature (°C) for Fe, Cu, Ni, Zn and Al Nitrates Thermal Decomposition

A reaction is possible at the standard conditions if its molar Gibb energy is negative. Therefore, according to the information contained on Figure (1), it can be observed that the decompositions occur at different temperatures. For example, while the copper nitrate thermal decomposition is possible for temperatures equal or higher than 200°C, the reaction associated with nickel nitrate starts at 600°C. For open systems, however, the thermodynamic tendency is defined in the same way by the sign of AG (AG = AG° + RTlnQ). In such systems, the reactions may occur at even lower temperatures, as a reaction quotient (RT.lnQ) can be provided, which is capable of making the total reaction Gibbs energy even more negative. It is known that this condition is relatively easy (operationally) to achieve when the equilibrium constant is close to 10"4. In these cases, although AG° > 0, it is possible to dissociate the nitrates with the continuous removal of gaseous products. On Table I some equilibrium constant values (Keq) were included for the five reaction systems under study.

Table I. Cu, Ni, Zn, Al and Fe Nitrates Thermal Decomposition equilibrium constants

T

CO 100 200 300 400 500 600 700

Keq

Fe

1.00E-10 1.00E+05 1.00E+21 1.00E+33 1.00E+42 1.00E+46 1.00E+50

Cu 5.30E-08 3.72E-01 2.08E+04 7.38E+07 4.63E+10 8.99E+12 7.57E+14

Ni 4.34E-44 5.60E-29 8.05E-19 1.88E-11 8.06E-06 2.42E-01 1.13E+03

Zn

4.92E-38 3.49E-24 7.21E-15 4.19E-08 6.38E-03 8.56E+01 2.09E+05

Al 1.19E-11 3.84E+01 1.72E+10 7.09E+16 1.80E+22 6.51E+26 5.64E+30

It can be noticed that equilibrium constants values between 10" and 10'4 are achieved at different temperatures for each nitrate. In praxis, as the decompositions are performed in open systems (gaseous products evolve freely), lower reaction temperatures can be applied.

180

Methodology

Controlled masses of hydrated nitrates (Fe, Cu, Ni, Zn, Al) were dissolved in deionized water, and the solution heated up to 400°C until all water vaporizes and the thermal decomposition of the precipitated nitrates settles in. This can be recognized through the evolution of gaseous N02 (a strong brown colored gas). After all NO2 escapes, the resulting oxide powder is weighted and characterized by x-ray diffraction and scanning electron microscopy. For all nitrates tested, the time elapsed from the beginning of the experiment to the end of the gas evolution was about to one hour.

Results and Discussion

Nitrates thermal decompositions The oxide conversion (a) achieved in each decomposition reaction was obtained through calculation of the expected oxide final mass (miheorctical), which depends on the reaction stoichiometry, and the experimental mass (m0btaincd) measured for each one of the decomposition products. The obtained data was included on Table II.

Table II. Oxide conversion achieved in each decomposition System

Fe

Cu Ni

Zn

Al

T(°C)

400

"> theoretical ( g )

2.055

2.145

2.001

2.020

2.001

""obtained (fi)

2.058 2.213

2.045

1.994

2.199

a ( % ) 99.8

96.8 97.8

98.7

90.1

It can be noticed that, in all cases, the oxide conversions are higher than 90% for the reaction temperature imposed. With the exception of the nitrates of Cu, Ni and Al, the observed deviations are lower than the expected experimental error (1.5%). This fact can be in a first glance explained by the presence of residual nitrates, which can be associated to the distinct kinetics behavior of each one of the decomposition reactions involved. For these nitrates, the heating of the obtained powders for extended times at 400°C would lead to a better agreement between the theoretical and obtained masses.

Materials characterization Figures 3 to 10 bellow present the results of the characterization of the decomposition products for each reaction system.

Fe reaction system On Figures 3, it is verified that Fe2Û3 (hematite) is the main reaction product of the iron nitrate thermal decomposition. Moreover, on Figure (4), it can be observed that the obtained oxide particles agglomerate into spherical clusters. The SEM image also suggests that the clusters are composed of nanostrucured crystals.

181

Figure 3. X-ray pattern of the Fe2Û3 sample

Figure 4. Fe203 sample, SEM, 800x

Cu reaction system From Figures 5 and 6, it can be noticed that CuO is prevailing oxide from the copper nitrate dissociation as well as the same tendency that appeared in Fe reaction system to form cristalline clusters with a spherical morphology. As in the case of Fe203, the CuO crystals present in the clusters have a size in the nanometric range.

Figure 5. X-ray pattern of the CuO sample

182

Figure 6. CuO sample, SEM, 800x

Figure 7. X-ray pattern of the NiO sample

Figure 8. NiO sample, SEM, 800x

183

Figures 7 and 8 show the formation of NiO as the main reaction product of the nickel nitrate pyrolisis. It is noticed that this oxide has a quite similar morphology to those identifed in Cu and Fe reaction systems. However, it can be observed that NiO cluster have a mean diameter higher than those other clusters, and as was verified for the oxides CuO and Fe203, the NiO clusters are also composed by nanosized particles.

Zn reaction system It is shown at Figure 9 that ZnO is major reaction product of zinc nitrate thermal decomposition. At Figure 10 it can be observed a distinct morphology as zinc oxide presents itself with a faceted crystals structure, completely different from those presented for the other oxides (Fe, Cu and Ni).

Figure 9. X-ray pattern of the ZnO sample

Figure 10. ZnO sample, SEM, 800x

Al reaction system Figures 11 and 12 confirm that AI2O3 is the resulting reaction product and is present as crystals in a pancake structure morphology. In the diffractogram (Figure 11), a strong contribution from an amorphous band can be identified. In the same signal, however, a clear broad peak around 66° is present, which indicates a significant contribution from the diffraction of nanosized crystals. The broadness of the mentioned peak indicates that this sample should have a much lower mean crystallite size as compared to the other oxides already discussed.

184

Figure 11. X-ray pattern of the AI2O3 sample

Figure 12. AI2O3 sample, SEM, 800x

Mean crystallite size Still regarding the materials characterization of the obtained oxides, a study of the x-ray diffraction patterns allows to calculate the mean crystallite size of the samples, as shown on Table III. It is interesting to observe that, in general, the crystallites of Fe2<D3, CuO and NiO are present with sizes within the same order of magnitude, between 35 and 60 nm. However, the crystallites of ZnO and AI2O3 are revealed, respectively, with larger and lower dimensions (157 and less than 1.0 nm) as compared to the oxides studied. The mean crystallite sizes, as determined from the quantitative evaluation of the diffractograms are consistent with the particle sizes observed in the SEM images. These results confirm the expectation that the proposed synthesis method allows the production of oxides containing nanosized particles.

Table 111. Mean crystallite size Oxide Sample

Fe203

CuO NiO

ZnO

AI2O3

Crystallite size (nm) 38.80 56.66 36.02 157.21 < 1.00

185

Final remarks

The thermodynamics considerations showed that it is possible to obtain oxides of Fe, Ni, Cu, Zn and Al from their respective nitrates thermal decompositions at fairly low temperatures (7"<600°C). In fact, these temperatures can be even lower once the dissociations happen are in praxis performed in open systems, which allow the evolution of gaseous products. The nitrates pyrolisis were conducted at 400°C, and enabled the achievement of high conversion levels (higher than 90%). The materials characterization though x-ray diffraction proved that Fe203, CuO, NiO, ZnO and AI2O3 were the prevailing decomposition products. Furthermore, scanning electron microscopy showed some distinct morphological aspects for each reaction product as Fe, Cu and Ni oxides tend to appear as particle clusters while the two others are formed either as faceted crystals (ZnO) or as crystals of pancake shape (AI2O3). Further analysis by the Rietveld method confirmed, through the mean crystallite size evaluation, that the samples contain particles with sizes in the nanometric range, in agreement with the SEM images.

References

1. M-S. Tsai, "Powder synthesis of nano grade cerium oxide via homogenous precipitation and its polishing performance", Materials Science and Engineering B, 110 (2004), 132-134

2. S. Kavecky, B. Janekovâ, J. Madejovâ, P. Sajgalik, "Silicon carbide powder synthesis by chemical vapour deposition from silane/acetylene reaction system", Journal of the European Ceramic Society, 20 (2000), 1939-1946

3. F. Danafar, A. Fakhru'1-Razi, M.A.M. Salleha, D.R.A. Biaka, "Fluidized bed catalytic chemical vapor deposition synthesis of carbon nanotubes—A review", Chemical Engineering Journal, 155 (2009), 37-48

4. X. Wu, Q-Z. Jiang, Z-F. Ma, M. Fu, W-F. Shangguan, "Synthesis of titania nanotubes by microwave irradiation", Solid State Communications, 136 (2005), 513-517

5. A. J. Reddy, M.K. Kokila, H. Nagabhushana, J.L. Rao, C. Shivakumara, B.M. Nagabhushana, R.P.S. Chakradhar, "Combustion synthesis, characterization and Raman studies of ZnO nanopowders", Spectrochimica Acta Part A: Molecular and Biomolecular Spectroscopy, 81 (2011)53-58

6. E.A. Brocchi, F.J. Moura, I.G. Solôrzano, M.S. Motta, "Synthesis and characterization of Cu and Ni bearing nano-composites and nano-structured alloys", Advanced Materials Research, 89-91 (2009), 65-72

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Fabrication of Lotus-type Porous Copper by Centrifugal Casting Technique

Yun-Soo Lee1, Hyeong-Tae Kim', Myoung-Gyun Kim2, Soong-Keun Hyun1

'School of Materials Science and Engineering, Inha University, Incheon, 402-751, South Korea 2Research Institute of Industrial Science and Technotogy(RlST), Pohang, 790-330, South Korea

Keywords: Lotus-type porous metal, Centrifugal casting. Directional solidification

Abstract

A centrifugal casting technique was used to fabricate lotus-type porous copper under hydrogen gas pressure similar to atmospheric. The molten copper was directionally solidified in the mould at a rotational velocity of 200-1000 rpm under a hydrogen gas pressure of 0,1 MPa, The results confirmed that porosity and pore size can be controlled by rotational velocity. Consequently, centrifugal casting is a promising fabrication process for safe and economical manufacture of lotus-type porous metals.

Introduction

Porous metals have lately attracted much attention in various industries because of their specific advantages such as light weight, fluid permeability, vibration and acoustic damping capacity, large surface area. Therefore, these metals are expected to be utilized as lightweight materials, vibration and acoustic energy damping materials, biomédical materials, and so on [1, 2],

Shapovalov [3] and Nakajima et al- [4] fabricated distinctive porous metals called gasar metals or lotus-type porous metals with cylindrical pores aligned in one direction. They have inherent characteristics by virtue of directional pores differ from those of conventional porous and foamed metals, such as superior mechanical properties, high thermal and electrical conductivity, smooth fluid permeability, and so on.

Conventional technique for fabrication of lotus metals have been studied by Hyun and Nakajima [5], Kashihara et al, [6] and Park et al. [7]; they fabricated lotus-type porous metals using the mould casting, continuous zone melting and continuous casting techniques at various conditions, respectively. Although these methods are convenience process, pressurized hydrogen gas is needed to control porosity and pore diameter.

In this study, we attempted to fabricate lotus-type porous copper by centrifugal casting at various rotational velocities under hydrogen gas pressure of 0.1 MPa. The effect of rotational velocity on porosity and pore diameter was also discussed.

Expérimenta) procedure

Lotus-type porous copper was fabricated by modified broken arm-type centrifugal casting apparatus, as illustrated in Figure 1. Pure copper was melted by induction coil in a graphite crucible under a hydrogen gas pressure of 0.1 MPa. As the melting temperature reached 1423 K, induction coil descend to bottom of chamber, and then molten copper is poured into the mould by centrifugal force at given rotational velocity. Lotus-type porous copper ingots were obtained by directional solidification in the mould. The rotational velocity was changed at 200, 300, 400, 500 and 1000 rpm, which was mechanically controlled by the electric motor.

The porosity (p) of solidified ingots was evaluated from following equation:

Apparant density of lotus copper I p(%)^\ 1-— — '-1—1x100

Density of nonporous copper J to

The apparent density of the individual ingot was measured by Archimedes' principle Then the ingots were cut in the directions parallel and perpendicular to the solidification direction by using a precision cut-off machine (Accutom-50; Struers GmbH, Copenhagen, Denmark). Each cross-section was polished with a series of emery papers and was observed using an optical microscope (VHX-200; KEYENCE Co., Japan). The pore diameter was measured in the cross-section perpendicular to the solidification direction using image analyzer (Image-Pro Plus; Media Cybernetics Co., USA).

Figure 1. Schematic drawing of the fabrication apparatus for lotus-type porous copper.

Results and discussion

Figure 2 shows the cross sectional views of the fabricated lotus copper perpendicular and parallel to the solidification direction. It was confirmed that the porosity and the average pore diameter decreased with an increase in rotational velocity but pore number density increased. It was much different compared to previous results fabricated by the conventional fabrication methods, because porosity and pore diameter are changed by pressurization or depressurization of hydrogen. It is reported that the porosity and pore diameter decrease with increasing hydrogen pressure by interaction of Sievert's law and Boyle's law [4], However, this research is performed under the same hydrogen gas pressure of 0.1 MPa, the amount of hydrogen dissolving in molten copper is

187

Figure 2. Cross sections perpendicular and parallel to the solidification of lotus-type porous copper fabricated at 200-1000 rpm under hydrogen gas pressure of 0.1 MPa. Pore diameter was obtained at cross section of fabricated porous copper up to 20 mm from bottom of ingot.

fixed. Accordingly, there is another mechanism different from conventional methods.

Figure 3 shows the distribution of pore diameter in the cross section up to 20 mm from the bottom of the mould in the lotus copper fabricated at various rotational velocities. It was confirmed that the frequency of fine pore increased with an increase in rotational velocity.

These results can be explained by Figure 4, which shows schematic drawing of the effect of rotational velocity on porosity and pore diameter during (a) pore nucleation and (b) growth. At the initial stage of centrifugal casting, melt convection and tubulence increase detached crystals [8]. These particles, in turn, act as heterogeneous nucleation sites for pores. At this time, hydrogen atoms dissolving in the molten copper diffused to the nearest nucleation site by diffusion field as illustrated in Figure 4-(a). However, the amount of hydrogen atoms is fixed according to Sievert's law, an increase in nucleation site results in pore refinement. At the same time, centrifugal force continuously put pressure on pores during pore growth as illustrate in Figure 4-(b). This force gradully increased with increasing rotational velocity. By the reason as above, the porosity and the average pore diameter decreased with an increase in rotational velocity but pore number density increased.

Figure 4. Mechanism of pore refinement during (a) pore nucleation and (b) growth. L: liquid, S: solid, G : gas pore.

Conclusion

Figure 3, Pore diameter distribution in lotus-type porous copper fabricated at various rotational velocities, (a) 200 rpm, (b) 300 rpm, (c) 400 rpm, (d) 500 rpm, (e) 1 OOOrpm.

A new process was developed for unpressurized fabrication of porous copper with directional pores. The porosity and average pore diameter decreased with increasing rotational velocity but pore number density increased. This is attributed to the amount of nucleation site and centrifugal force by changing the rotational velocity.

188

Acknow ledgements

This research was supported by the Fundamental R&D Programs for Core Technology of Materials funded by Ministry of Knowledge Economy Korea

References

[1] A.G. Evans, J.W. Hutchinson and M.F. Ashby, "Multifunctional ity of cellular metal systems," Progress in Materials Science, 43 (1999), 171-211. [2] J. Banhart, "Manufacture, characterization and application of cellular metals and metal foams," Progress in Materials Science, 46(2001), 559-632. [3] V. Shapovalov and L. Boyko, "Gasar-A New Class of Porous Materials," Advanced Engineering Materials, 6(2004), 407-410. [4] H. Nakajima, "Fabrication, properties and application of porous metals with directional pores," Progress in Materials Science, 52 (2007), 1091-1173. [5] S.K. Hyun and H. Nakajima, "Effect of solidification velocity on pore morphology of lotus-type porous copper fabricated by unidirectional solidification," Materials Utters, 57 (2003), 3149-3154. [6] M. Kashihara, S.K. Hyun, H. Yonetani, T. Kobi and H. Nakajima, "Fabrication of lotus-type porous carbon steel by unidirectional solidification in nitrogen atmosphere," Scripta Materialia, 54 (2006) 509-512. [7] J.S. Park, S.K. Hyun, S. Suzuki and H. Nakajima, "Effect of transference velocity and hydrogen pressure on porosity and pore morphology of lotus-type porous copper fabricated by a continuous casting technique," Ada Materialia, 55 (2007) 5646-5654. [8] S.R. Chang, J.M. Kim and C.P. Hong, "Numerical Simulation of Microstructure Evolution of Al Alloys in Centrifugal Casting," ISIJInternational, 41 (2001) 738-747.

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EFFECTS OF PULSED MAGNETIC ANNEALING ON THE GRAIN BOUNDARY OF PRIMARY RECRYSTALLIZED MICROSTRUCTURE IN

THE GRAIN-ORIENTED SILICON STEEL

Junjun Huang, Lihua Liu, Xin Xia, Xiang Jiang, Lijuan Li *, Qijie Zhai

School of Material Science and Engineering, Shanghai University Shanghai, 200072, China, Email: esneullj@163.com

Keywords: Pulsed Magnetic Field; Annealing; Recrystallization And Grain Boundaries; Grain Oriented Silicon Steel; EBSD Technology

Abstract

In this work, the effects of pulsed magnetic field applied during the annealing of grain-oriented silicon steel on the grain boundaries in the primary recrystallizated microstructure were investigated. Samples of cold rolled grain-oriented silicon steel were annealed under pulsed magnetic field with the maximum strength IT from three different directions-rolling direction, transverse direction and normal direction at the temperature of 700°C for 16 minutes. Electron Backscattering Scanning Diffraction (EBSD) technology was used to measure grains for texture determination, and software called Channel 5 was used to calculate the messages of boundaries. Results show that pulsed magnetic field can influence the development of grain boundaries. It is found that the frequency of the low angle boundaries increases when pulsed magnetic field is applied, especially from the rolling direction. Compared to the ordinarily annealed sample, the frequencies of CSL boundaries vary in samples annealed with pulsed magnetic field in different directions.

Introduction

The grain-oriented silicon steel is widely used in transformers, motors and power generators. L.K. Fionova [1] pointed that the GB plane, density of GB defects, GB spatial distribution, i.e. mutual location of various kinds of GB, GB inclination (grain shape) and GB density (grain size) or size of clusters of low energy GBs, can be used for GB engineering. In order to improve its magnetic properties for energy saving by means of GB engineering, DC magnetic field is applied to control the micro-structure through the grain growth and texture development which are closely related to the grain boundaries [2-6]. In those researches, the DC magnetic field is found to greatly influence the development of boundaries. Other kind of magnetic field is rarely applied for researching in grain boundary related phenomena. H.PENDER [7] researched the behavior during the annealing of steel in an alternating magnetic Field in 1913. Recently Harada K. et al. inspected the effect of DC and AC magnetic field on grain growth in electrodeposited nanocrystalline nickel which indicated that the application of a magnetic field could enhance grain growth, but differently at the very early stage and the late rapid grain growth stage of annealing due to the change of the grain boundary mobility [8]. However, the study of pulsed magnetic field applied during annealing of deformed materials is randomly published. As generally considered, pulsed magnetic field can cause the soft magnetic materials suffer repeated dynamic magnetostriction which will significantly affect the mobility of grain boundaries. So this paper is to focus attention on the effects of pulsed magnetic annealing on the formation of grain boundary in primary recrystallized microstructure with the grain-oriented silicon as the experimental materials.

191

Experimental Procedures

The experimental materials in this study were industrially 0.3mm thickness cold rolled grain-oriented silicon steel strips after 87% reduction. The ranges of chemical compositions in the starting materials are listed in Table I , and A1N, MnS are employed as inhibitors for this silicon steel.

Table I . Chemical compositions of the starting materials

Elements

Contents in wt%

C

0.06-0.08

Si

3-3.3

N

0.007-0.009

Al

0.02-0.03

Fe

Bal.

Pulsed magnetic field is formed by a columned solenoid of die diameter O=70mm with power supplied by a self made high voltage pulsed power resource. A range of 60mm length along the axial direction of the solenoid is of homogenous pulsed magnetic field. This pulsed magnetic field direction is parallel to the axial direction. The strength of pulsed magnetic field increases from zero to the crest (lTesla) in the approximately sinusoid and then attenuates in the exponent curve. Generally viewed, the half width of pulsed magnetic field approximates to 50ms and its frequency is 0.5Hz. A tube furnace is in the solenoid with its center positioned in the center of the solenoid. A length of 80mm along the axial direction is of homogeneous temperature in the middle of the furnace, whereas the temperature at the end of die furnace viz. outside of the solenoid approximates to the room temperature. The furnace is of dynamic seal, and a stainless pole of which the end can be moved into the middle of the furnace or outside of the solenoid.

Specimens can be positioned at the end of the drawing and pushing role in three ways: pulsed magnetic field direction parallels the rolling direction, or transverse direction, or normal direction respectively. The specimens were firstly positioned at the end of the pole outside of the solenoid, and then the furnace was sealed and vacuumed, and then filled wiüi high pure argon. When the temperature in the middle of furnace reached the setting temperature (700°Q, specimens at die end of the pole were pushed into the middle of the furnace wiüi the pole, and timing began at the same time. After 16 minutes, specimens were pulled out to stay in the cool area of the furnace outside of the solenoid for air-cooling. The specimen annealed without magnetic field is denominated OS, and denoted by 700-16-0, while the specimen magnetic annealed in the direction of rolling direction (RD) is denominated MR, and denoted by 700-16-1T-RD. Similarly the specimen magnetic annealed in the direction of transverse direction (TD) or normal direction (ND) is denominated MT or MN, and denoted by 700-16-1T-TD or 700-16-1T-ND.

Specimens were grinded and electric polished to the subsurface layers, and then grain textures were detected with EBSD technology at a magnification of 100 times by the scanning electric microscopic named APOLLO 300. The scanning area covered 750um (transverse direction) x 300um (rolling direction), and the scanning step is 1.5um. The software Channel-5 from Oxford Instrument was used to obtain information about grain boundaries.

192

Results and Discussion

Figure 1. The distribution of misorientation angles in the four samples

Figure 1 illustrates the distribution of misorientation angles in the four samples including the low energy boundaries (misorientation angles <15°), the high energy boundaries (misorientation angles are between 20° and 45°) and the remained boundaries. The behavior of the grain boundary migration occurs through the exchange of vacancies, and the high energy boundary more mobile for it is more capable of absorbing and emitting vacancies than the low-energy boundary [9],

In contrast with that in OS, the frequencies of the low energy boundary increase in pulsed magnetic annealed samples, especially in MR of which the frequency of low energy boundaries significantly increases by about 69%. However to the high angle boundaries, the frequencies in magnetic annealed samples decrease to some extent. That means pulsed magnetic field during annealing of grain-oriented silicon steel favors the occurrence of low energy grain boundaries in the recrystallized microstructure, which is similar to that with the statistic magnetic field [6,10],

The distribution of low I CSL interfaces (i.e. £3-29) in the four samples is illustrated by Figure 2, which shows £3 is the main component, and the frequency of twin boundary (13) in MR increases while decrease in MT and MN.

Some classifications are made for comparisons, and the summarized details are listed in the below Table II.

193

Figure 2. The distribution of the low £ CSL interfaces in the four samples

Table II. Comparisons of all kinds of classifications of CSL boundaries in the four samples

£3-29 £3-49

23 25-49

25+27+29 2<100> 2<U0> 2<111>

Low 2 CSL of high energy boundaries

OS 9.18 12.04 3.16 8.88 1.88 0.92 2.04 4.52 3.33

MR 9.49 12.23 3.47 8.76 1.55 0.71 2.02 4.71 2.86

MT 8.39 11.22 2.90 8.32 1.68 0.77 1.81 4.06 2.93

MN 8.37 11.33 2.93 8.4 1.70 0.74 1.91 4.02 2.94

Compared to OS showed in Table II, the total frequency of low I CSL interfaces (£3-29) in MR increase, while decrease almost the same in MT and MN. Since £3 has low mobility, the frequencies of mobile CSL boundaries (Z5-29) in the three pulsed magnetic annealed samples decrease compared with that in OS.

According to the CSL theory, the CSL boundaries namely 15 [11], X7 [12], 19 [13] are responsible for the growth of Goss grains. The total frequency of 15, 17 and £9 which are responsible to the growth of Goss grains in CSL theory decreases when pulsed magnetic is applied during annealing, and it decreases the most in MR. That means in the incipient recrystallization, pulsed magnetic field makes Goss grains more stable.

The CSL boundaries are often described by defining the rotation axis and the rotation angle, so here I <uvw> denotes the CSL boundaries with the rotation axis <uvw>. I<100> includes I 5,13a, 17a, 25a, 29a, 37a, 41a. I<110> includes 19, 11, 19a, 27a, 33a, 33c, 41c. 2<111>

194

includes I 3, 7, 13b, 19b, 21a, 31a, 37c, 39a, 43a, 49a. Table 2 shows that the frequencies of these three kinds of CSL boundaries in pulsed magnetic annealed samples are decreased correspondingly compared with that in OS except that the frequency of I<111> in MR is higher than in OS.

During those low I CSL interfaces, some are also the high energy boundaries, like 5, 7, 9, 13a, 13b, 17a, 19a, 21a, 21b, 23, 27a, 27b and 29a. Comparing the summarized numbers, these numbers of pulsed magnetic annealed samples are fewer than that in OS.

In inclusion, pulsed magnetic field makes the recrystallized grains in MR, MT and MN more stable than that in OS.

As the main component, gamma texture is analyzed in details as below. Figure 3 illustrates the OIM map of gamma textures component surrounding by all kinds of CSL interfaces and misorientation angle boundaries in the four samples.

Figure 3. The OIM map of gamma textures with all CSL interfaces and low (thin black curve) and high energy boundaries (wide black curve) around it in the four samples (a) OS, (b) MR, (c)

MT, (d) MN.

Table III. The frequencies of low and high misorientation angles boundaries around all grains and gamma texture in the four samples

Frequency of boundaries /% Around all grams

Around gamma texture

Misorientation angles

<15° 20-45° <15°

20-45°

700-16-OT

0.09 0.47 0.17

0.43

700-16-1T-RD

0.15 0.43 0.33

0.31

700-16-1T-TD

0.14 0.46 0.35

0.30

700-16-1T-ND

0.13 0.46 0.39

0.27

Table III shows that the gamma oriented grains have a higher number of small misorientation angle boundaries (< 15°) than the specimen as a whole. Pulsed magnetic field enhances the

195

development of small misorientation angle boundaries especially surrounding gamma texture. Magnetic field applied from the normal direction can most effectively promote the development of small misorientation angle boundaries around gamma grains. The difference of grain boundary distributions around all grains respectively in MR, MT and MN is quite small.

Figure 4. The distribution of low ICSL interfaces around gamma texture in the four samples

The total frequencies of low I CSL interfaces around gamma texture in magnetic annealed samples are smaller than that of OS. The frequency of S 3 in MT decreases the most compared to that in OS, while the frequencies of I 7 in magnetic annealed samples are much lower than in OS.

In order to explain these changes, some theories are proposed. Pulsed magnetic field promotes the recrystallized grains to reach the relative stable «crystallization stage due to the effects below: the vibration effect which makes stress relieve, the improvement in mobility of dislocation which helps stress pinning the dislocation relieved and the relax of interaction of dislocation with domain wall [14]. So magnetic field shortens the time for recovery. And then during the following recrystallization, magnetostrictions in MT and MN are larger than in MR [15]. Therefore the grains in MT or MN quickly develop almost at the same time, which makes recrystallization in MT or MN reach a more stable stage with small average grain size. The corresponding evidence is that the average grain sizes measured by EBSD technology of the four samples are 8.63um, 8.46um, 7.84um and 8.34um in OS, RD, TD and ND respectively. And when the grain size decrease, the frequency of low energy boundaries like low angle boundaries and low Z CSL boundaries increases [10]. That well explains why the frequency of low angle boundaries increases in magnetic annealed samples.

Conclusions

It is found that the frequency of the low angle boundaries increases when pulsed magnetic field is applied, especially from the rolling direction. Compared with that in OS, the total frequency of low £ CSL boundaries (13-29) increases in MR, while decrease almost the same level in MT and MN. Moreover, the total frequency of Z5,17 and 19 which are responsible to the growth of Goss

196

grains in CSL theory decreases when pulsed magnetic is applied, and it decreases the most in MR. The total frequencies of low S CSL interfaces around gamma texture in magnetic annealed samples are lower than that of OS. Pulsed magnetic field applied from the normal direction can most effectively promote the development of small misorientation angle boundaries around gamma grains. All in all, pulsed magnetic field makes the recrystallized grains in MR, MT and MN quickly become more stable than in OS. The development of grain boundaries can be controlled by applying pulsed magnetic field during annealing of cold rolled silicon steel.

Acknowledgement

This project is supported by the National Nature Science Foundation of China (Grant No.50804028). The authors wish to express their thanks to Instrumental Analysis and Research Center of Shanghai University for EBSD technology.

References

1. L.K. Fionova, Y.A. Lisovskii, and T. Watanabe, "Grain Boundary Design and Control for Electrical and Properties of Polycrystalline Materials," Materials Letters, 28 (1996), 141-147.

2. C.M.B. Bacaltchuk et al., "Effects of Magnetic Field Applied During Secondary Annealing on Texture and Grain Size of Silicon Steel," Scripta Materialia, 48 (9) (2003), 1343-1347.

3. C.M.B. Bacaltchuk et al., "High Magnetic Field Effect on Texture and Grain Growth of GNO Silicon Steel," Materialwissenschaft und Werkstofftechnik, 36 (10) (2005), 561-565.

4. E. Liu et al., "Primary Recrystallization of Grain Oriented Silicon Steel Strip Rolled by CSR and Annealed in Magnetic Field," Journal of Material Science and Technology, 21 (4) (2005), 455-458.

5. Y. Xu et al., "Effects of Strong Magnetic Field on Recrystallization and Coarsening of Grains During Annealing in 3% Silicon Steel," Journal of Magnetics Society of Japan, 24 (4-2) (2000), 651-654.

6. N. Masahashi, M. Matsuo, and K. Watanabe, "Development of Preferred Orientation in Annealing of Fe-3.25% Si in Hign Magnetic Field," Jounal of Material Research, 13 (2) (1998), 457-461.

7. H. Pender and R.L. Jones, "The Annealing of Steel in An Alternating Magnetic Field," Physical Review Series II, 1 (4) (1913), 259-273.

8. K. Harada et al., "Enhancement of Homogeneity of Grain Boundary Microstructure by Magnetic Annealing of Electrodeposited Nanocrystalline Nickel," Scripta Materialia, 49 (2003), 367-372.

9. R.Z. Valiev et al., "Grain Boundary Structure and Properties Under External Influences," Physica Status Solidi (a), 97 (1) (1986), 11-56.

10. T. Watanabe et al., "The Effects of Magnetic Annealing on Recrystallization and Grain-boundary Character Distribution (GBCD) in Iron-cobalt Alloy Polycrystals," Philosophical Magazine Letters, 62 (1) (1990), 9-17.

l l .J . Harase, R. Shimizu and D.J. Dingley, "Texture Evolution in the Presence of Precipitates in Fe-3% Si Alloy," Ada Metallurgica et Materialia, 39 (5) (1991), 763-770.

12. J. Harase, "Texture Evolution in the Presence of Precipitates in Fe-3% Si Alloy," Materials Science Forum, 94-96 (1992), 419-424.

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13. Y. Yoshitomi et al., "Influence Growth of Primary Recrystallized Grains on Secondary Recrystallization Texture in Fe-3% Si Alloy," Materials Science Forum, 113-115 (1993), 715-720.

14. S.N. Prasad, P.N. Singh and V. Singh, "Influence of Pulsating Magnetic Field on Softening Behaviour of Cold Rolled AISI 4340 Steel at Room Temperature," Scripta Materialia, 34 (12) (1996), 1857-1860.

15. H.<I>. Dubulof, H.H. Lapkin, ed. S. Peng, Electrical Steel (Beijing, BJ: China Industry Press, 1965), 62.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

RELATIONSHIP BETWEEN HEAT INPUT AND MICROSTRUCTURE AND MECHANICAL PROPERTIES OF LASER BEAM WELDED

SUPERALLOY INCONEL 718

Akin ODABASI1, Necip ÜNLÜ2, Gültekin GÖLLER2, M. Niyazi ERUSLU2, E. Sabri KAYALI2

'Firat University, Faculty of Engineering, Department of Metallurgical and Materials Engineering,

23119, Elazig, Turkey

Istanbul Technical University, Faculty of Chemistry-Metallurgy, Materials Science andMetallurgical Engineering Department,

34469, Maslak, Istanbul, Turkey

Keywords: Inconel 718, laser welding, solidification rate, hardness

Abstract

Autogenous butt, laser beam welds were carried out using Inconel 718 alloy sheets (2.1 mm thick). The relationship between heat input laser beam welding (LBW) and the microstructural and mechanical properties of superalloy Inconel 718 were investigated. Nine different heat inputs in the range of 61.29-90.09 Jmm"' were applied to evaluate the geometry of weld seams. Full penetration was achieved in all weld experiments. Optical and field emission scanning electron microscopy and microhardness tests were performed. Test results indicate that increasing the amount of heat input from 61.29 Jmm"1 to 90.09 Jmm'1 presented solidification rates between 1.27xl05 and 1.46xl06°Cs"'. The hardness property of the weld samples decreased with the increasing the heat input.

Introduction

Using the laser beam welding (LBW) in several industries such as automotive, aerospace, military and power plants has been increased continuously due to the unique properties i.e., rapid welding speed, low heat input, narrow fusion zone, and heat-affected zone (HAZ) [1, 2], Since 1990s, the weldability, high resistance response to weld solidification cracking, excellent mechanical properties, and oxidation resistance at elevated temperatures [2-4] have been caused a significant interest in a Ni-based superalloy Inconel 718. Although various investigations on high-energy-density beam welding processes such as electron beam (EB) and gas tungsten arc (GTA) welds [3, 4] have been carried out, there have been a limited number of literature on LBW of superalloy Inconel 718 [2, 5, 6,]. The goal of this investigation was to determine the relationship between the heat input and microstructure and mechanical properties of laser beam welded superalloy Inconel 718.

Experimental procedures

Autogenous butt, laser beam welds were performed using Inconel 718 alloy sheets (2.1 mm thick). Table 1 and Table 2 summarize the composition of the base metal and the welding parameters, respectively. Argon was used as the shielding gas. The sections tranverse to the welding direction were cut by abrasive cutter (SamplMet 2). Then the sections were prepared using standard metallographies techniques and etched with the solution of Glyceregia (3 parts

199

glycerol, 3 parts HC1 and 1 parts HNO3). The microstructures of the welded samples were investigated using an optical microscope model of Leica DMRX and a field emission scanning electron microscopy (FESEM) JEOL JSM 7000F with fully computer controlled with EDS, WDS and EBSD attachments. Microhardness measurements were performed using a Vickers microhardness tester of Leica VMHT MOT at a load of 200 g.

Table 1. Composition (in wt %) of the base material. Elements

SAE |7|

Measured

C

0-0.08

Mn

0-0.35

Si

0-0.35

0.32

P

0-0.015

0.009

S

0-0.015

Cr

17-21

17.2

Ni

50-55

55,4

Mo

2.8

3.3

2.8

Nb

4,75

5,50

4.33

Ti

0.65

1.15

0.7

Al

0.2

0.8

0.3

Co

0-1

0.13

Ta

0-0.05

B

0-0.006

-

Cu

0-0.3

0.02

Fe

Remain

18.96

Table 2. Welding parameters.

Sample No#

1 2 3 4 5 6 7 8 9

Laser Power (Watt)

2850 2850 2850 2500 2500 2500 2000 2000 2000

Welding Speed (mm/s)

46.5 36.4 31.1 36.4 31.1 26.4 31.1 26.4 22.2

Plasma Suppression and Inert Gas Application (1/min)

Plasma

30

Trail

15

Back

10

Heat Input (J/mm)

61.29 78.29 91.63 68.68 80.38 94.69 64.30 75.75 90.09

Results and discussion

The effect of heat input on the microstructures of laser beam welded Inconel 718

Figure 1 shows the effect of laser beam power on the geometry of weld seams based on the welding speed. At the constant laser beam power such as 2000 watt, increasing the welding speed from 22.2 mm/s to 31.1 mm/s changed the resulting weld shape from a hourglass shape to a nail shape. Similar changing was observed at the constant welding speed i.e. 31.1 mm/s while decreasing the laser beam power from 2850 watt to 2000 watt. The variation of the seam widths as a function of the heat input is determined. Increasing the heat input from 61.29 J/mm to 90.09 J/mm exhibits significant increase in seam widths, i.e., the seam widths of the mid-region, and the root region are 686 urn, and 1053 urn vs., 1097 um and 1677 urn, respectively. These results indicate that both the heat input and the welding speed have significant effect on the welding geometry. The effect of heat inputs on the weld seam microstructures were determined by using optical microscope (OM) investigations. Figure 2 (a-c) shows the the solidification microstructures based on applied heat inputs. When the heat input increased from 61.29 J/mm to 90.09 J/mm, the measured average dendrite-arm spacing was linearly increased from 1.49 urn (Fig. 2a.) to 2.59 um (Fig. 2c). According to Wang et al. [8], Bouse and Mihalisin [9], and Won et al. [10], the solidification rate was calculated. Figure 3 shows the relationship between the heat input and the solidification rate. The calculated solidification rates are consistent with the weld metal solidification rates [2, 11],

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Figure 1. The effect of laser beam power on the geometry of the weld seams based on the welding speed.

Figure 2. The effect of the applied heat inputs on the weld seam microstructures obtained at the heat inputs of (a) 61.29 J/mm, (b) 80.38 J/mm, and (c) 90.09 J/mm.

201

Figure 3. The relationship between the heat input and the solidification rates calculated as a function of dendrite arm-spacing according to Wang et al. [8], Bouse and Mihalisin [9], and Won et al. [10].

Figure 4. The SEM micrographs taken from the weld fusion zones of the welded samples obtained at the heat inputs of (a) 61.29 J/mm, (b) 80.38 J/mm, and (c) 90.09 J/mm.

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Laves phase is identified as A2B (Ni, Fe, Cr)2(Nb, Mo, Ti, Si) form, and its formation requires a niobium concentration ranging from 10 pet to 30 pct.[12-14]. It has been also reported that the amount of the Laves phase and its Nb concentration are a function of solidification conditions. High cooling rates in laser beam welding indicate a lesser extent of Nb segregation due to the insufficient times for solute distribution. To evaluate the fusion structure and determine the phases in detail, the SEM investigation was performed. The SEM micrographs taken from the weld fusion zones of the welded samples obtained at the heat inputs of 61.29 J/mm, 80.38 J/mm and 90.09 J/mm are shown in Figures 4(a), (b) and (c), respectively. The EDS spectra results [15] taken from the structure labeled as L in Figure 4 (a) and (c) conformed with the Laves phase composition. The current Nb concentrations observed in samples #1 and #9 are 21.27 wt % and 26.13 wt %, respectively [15]. These data conform with the findings of Ram et al.[4], Jawwad et al. [16], and Radhakrishna et al.,[17]. The amount of Nb in Laves phase of EB welds of Inconel 718, about 19.62 wt pet and 38.34 wt pet, were estimated by Jawwad et al. [16] and Vincent [18] respectively, indicating the wide range of Nb concentration in the Laves phase and the effect of weld metal cooling rate.

Microhardness investigation

Figure 5 and Figure 6 show the hardness profiles of the welds obtained using both low and high heat inputs, respectively. The hardness range of the BM used in this study was around 240-250 HV0.2. The average hardness values of the FZ areas of the welds at the heat inputs of 61.29 J/mm, 80.38 J/mm, and 90.09 J/mm, were decreased to approximately 330 HV0.2, 300 HV0.2, and 280 HV0.2, respectively. Decreasing the hardness property has been attributed to the higher niobium segregation, coarser Laves morphology, and a higher amount of Laves phase because of the low weld metal cooling rates [6].

Figure 5. Microhardness profile of the weld fusion zone, heat affected zone and base metal for the Sample # 1.

203

Figure 6. Microhardness profile of the weld fusion zone, heat affected zone and base metal for the Sample # 9.

Conclusions

The effect of laser beam power on the geometry of weld seams, and microstructure and hardness properties of Inconel 718 superalloy having equaxied austenitic structure was investigated using nine different heat inputs in the range of 61.29 - 90.09 J/mm. Increasing the heat input resulted with a nail shape welding bead geometry. Microstructural investigation exhibited fine dendritic structure of the weld fusion zone obtained with the low heat input (61.29 J/mm) due to the rapid weld metal cooling rates. When the heat input increased from 61.29 J/mm to 90.09 J/mm the measured average dendrite-arm spacing was linearly increased from 1.49 urn to 2.59 um. Further, observing coarser Laves phase verified the coarse dendritic welding structures and hardness property at the low cooling rates.

Acknowledgement

This work was financially supported by the Istanbul Technical University, Turkey, Research Fund (Project #32101). The authors wish to thank Mr. Hüseyin Sezer and Mr. Tamer Alpak for their help and contributions during the SEM analysis of this study. Also, the authors thank to Mr. Hüseyin Gökcan, Mr. Kubilay Yildirim (GE Marmara Technology Center), Mrs. Ay§e Hilal Ulukardesjer and Mr. Mizrap Cambeyaz for their help.

References

1. G. Cam and M. Koçak, "Progress in Joining of Advanced Materials," International Materials Reviews, 43, 1 (1998) 1-44. 2. G.D. Janaki Ram et al., "Microstructure and Tensile Properties of Inconel 718 Pulsed Nd-YAG Laser Welds," Journal of Materials Processing Technology, 167 (2005), 73-82. 3. C.V.S. Murthy, "Electron beam welding of alloy 718 - a study on the effects of beam oscillation techniques" (M.Sc. Thesis, Department of Metallurgical Engineering, Indian Institute of Technology, Madras, Chennai, India, Jan 2004).

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4. G.D.J. Ram et al., "Control of laves phase in Inconel 718 GTA welds with current pulsing," Science and Technology of Welding and Joining, 9, 5 (2004), 390-398. 5. J.K. Hong et al., "Microstructures and Mechanical Properties of Inconel 718 welds by CO2 Laser Welding" Journal of Materials Processing Technology, 201 (2008),515-520. 6. G.D.J. Ram et al., "Microstructure and Mechanical Properties of Inconel 718 Electron Beam Welds," Materials Science and Technology, Vol 21, 10 (2005), 1132-1138. 7. SAE AMS 5596J, Aerospace Material Specification, Nickel Alloy, Corrosion and Heat Resistant, Sheet, Strip, Foil and Plate, 52.5Ni-19Cr-3.0Mo-5.1Cb-0.9Ti-0.5Al-18Fe Consumable Electrode or Vacuum Induction Melted 1775 °F (968 °C) Solution Heat Treated, Rev. July 1999, SAE International. 8. H.M. Wang et al., "Rapidly solidified MC carbide morphologies of a laser-glazed single-crystal nickel-base superalloy" Materials Science and Engineering A, 156, 1, (1992), 109-116. 9. G.K. Bouse and J.R. Mihalisin: in Superalloys, Supercomposites and Superceramics, J.K. Tien and T. Caulfield, eds., Academic Press, Burlington, MA, (1989), 99-148. 10. Y.M. Won et al., "Effect of cooling rate on ZTS, LIT and ZDT of carbon steels near melting point" 1SIJInternational, 38, 10, (1998), 1093-1099. 11. S.A. David and J.M. Vitek, "Principles of Weld Metal Solidification and Microstructures," Conference Proceedings on Trends in Welding Research, ed. S.A. David, J.M. Vitek, Gatlinburg, TN, USA, June 1-5, 1992, 147. 12. M.J. Cieslak et al., "A comparison of the solidification behavior of Incoloy 909 and Inconel 718" Metallurgical and Materials Transactions A, 1990, 21A, 1, (1990), 479-488. 13. C. Radhakrishna and K.P. Rao, "The formation and control of Laves phase in superalloy 718 welds" Journal of Materials Science, 32, 8, (1997), 1977-1984. 14. J.J. Schirra, R.H. Caless, and R.W. Hatala, "The effect of laves phase on the mechanical properties of wrought and cast + HIP Inconel 718" Proc. Conf. Superalloy 718, 625 and Various Derivatives, E.A. Loria, ed., TMS, Warrendale, PA, (1991), 375-388. 15. A. Odabaçi, "Optimization of Parameters for Carbon Dioxide Laser Welding of Superalloys" (M.Sc. Thesis, Istanbul Technical University, Institute of Science and Technology, 2010, in Turkish). 16. A.K.A. Jawwad, M. Strangwood, and C.L. Davis, "Microstructural modification in full penetration and partial penetration electron beam welds in INCONEL-718 (IN-718) and its effect on fatigue crack initiation" Metallurgical and Materials Transactions A, 36A, 5, (2005), 1237-1247., 17. C. Radhakrishna, K.P. Rao, and S. Srinivas, "Laves phase in superalloy 718 weld metals" Journal of Materials Science Letters, 14,24,(1995), 1810-1812. 18. R. Vincent, "Precipitation Around Welds in the Nickel-Base Superalloy, Inconel 718" Acta Metallurgica, 33, 7, (1985), 1205-1216.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EXPERIMENTAL STUDY ON THE BEHAVIOR OF SLAG ENTRAPMENT AND INCLUSION REMOVAL IN 44 T LADLE WITH

ARGON BLOWING

Shu-guo Zheng, Miao-yong Zhu, Zhong-fu Cheng

School of Materials & Metallurgy, Northeastern University; 3-11 Wen Hua Rd.; Shenyang, Liaoning, 110819, P. R. China

Keywords: Ladle with argon blowing, Powder entrapment, Inclusion behavior, Water modeling

Abstract

The powder entrapment and inclusion behavior in 44 ton ladle with argon blowing was studied in a water model. The effect of gas flowrate on powder entrapment in the ladle was investigated by choosing the mixture of kerosene and vacuum pump oil simulated as slag, and the effects of time and gas flowrate on inclusion removal were investigated by choosing emulsion drops simulated as inclusions. The results show that the critical gas flowrate for powder entrapment was 0.66 Nl/min, and the powder entrapment occurred continuously with those gas flowrates larger than 3.31 Nl/min. Most of the inclusions can be removed in eight minutes, and all the inclusions which can be removed almost disappeared from the system in twenty-eight minutes. The inclusion removal rate decreased first and then increased with the increase of gas flowrate. Furthermore, it had been found that the inclusion removal was in exponential relationship with gas blowing time.

Introduction

The secondary refining in the ladle has received a considerable attention in the production of high quality of steel. And gas stirring in a ladle is one of the most common refining techniques. Besides homogenizing the temperature and chemical composition, another function of gas stirring is to reduce the inclusions in molten steel as much as possible. In the past decades, many works had been done in respect of the fluid and mixing in gas-stirred ladles [1-4], while just several investigations concerned the inclusion behaviors. Among those studies on inclusion behavior, the analysis of steel samples [5,6] and mathematical modeling [7-9] is more than physical modeling [10,11], As slag entrapment has a bad effect on the steel quality, it should be avoided in the gas-stirred ladle. In recent years, a new physical simulation method that the collision and aggregation of inclusions in molten steel can be modeled had been gradually applied to study the inclusion behavior during the secondary processes [10-12]. With the same physical modeling method, based on the investigation of behavior of slag entrapment in a 44 ton gas-stirred ladle, the effects of time and gas flowrate on inclusion removal were investigated in a water model. And the process of inclusion removal was also studied.

Experimental

Similarity Considerations

A 1:4 down-scale physical model that simulated a 44 ton gas-stirred ladle was used. The main parameters for model and prototype are listed in Table I.

207

Table I. Main parameters for model and prototype

Prototype Model

Ladle top diameter (mm) 2168 542

Ladle bottom diameter (mm) 1821 455

Ladle depth (mm) 2100 525

Water was chosen to simulate the molten steel and air was chosen to simulate argon gas in actual system. According to the similarity rule, the physical model should have a same Froude number as the prototype [13], therefore the gas flowrates in model corresponding to those in prototype can be calculated. The model slag was the mixture of kerosene and vacuum pump oil, and the proportion of them satisfied the principle that the kinematic viscosity ratio of the mixture oil to water equals to that of slag to liquid steel [11]. A certain emulsion drops were used to simulate the inclusions, and the criteria for choosing the model inclusion as well as the relationship between the water model and prototype for inclusion size had been made detailed description elsewhere [10-12]. The relevant parameters of the media for model and prototype are shown in Table II, and the relationship for inclusion size between the ladle model and prototype can be obtained. The emulsion drops with the equivalent diameter of 376.7 um in model ladle can simulate A1203 inclusions with 97.1 urn or SiC>2 inclusions with 83.0 um in equivalent diameter in molten steel.

Table II. Relevant parameters of the media for model and prototype

Liquid density (103kg-m"J) Inclusion density (103kg-m'3)

Gas density* (103kg-m"3) Slag kinematic viscosity (m -s"')

Gas flowrate (Nl-min"1)

Prototype 7.1 (Steel)

3.9(AI203)or2.7(Si02) 1.784 (Ar) 3.8xl0-6

10-300

Model 1.0 (Water solution) 0.985 (Emulsion)

1.205 (Air) 4.2x10-6

0.22-6.63 * 7=273 K,p=1.0xl05 Pa

Experimental Method

The effect of gas flowrate on powder entrapment in the ladle was first investigated and the critical gas flowrate for powder entrapment was obtained. Considering the bad effect of slag entrapment on the steel quality, under the condition of avoiding powder entrapment, the effects of time and gas flowrate on inclusion removal were investigated. The schematic diagram of the experimental setup is shown in Figure 1.

In the study on the behavior of slag entrapment, air was blown into the model ladle through a porous plug located in the model ladle bottom. The model slag with the thickness of 20 mm was added by a glass rod and distributed homogenously on the model ladle surface before starting gas blowing. The behavior of slag entrapment was observed and recorded by the naked eye under different gas flowrates and the slag entrapment phenomena was recorded by a camera.

In the investigation of inclusion removal, air was blown into the model ladle through a porous plug. 30ml of emulsion was added into the model ladle and distributed homogenously before starting gas blowing. Water inflows from the bottom of model ladle with a very low flowrate only to realize an overflow condition to removal the floated inclusions from the top free surface. The mixture of emulsion and water solution was collected at a fixed time interval. The pure emulsion can be separated by dealing with the collected mixture, and the inclusion removal rate at a certain time can be given by the following equation:

208

XX

where rjt is the sum of removal rate from the first time interval to the kth time interval, and Va

and V, are the initial emulsion volume in the model ladle and the removal emulsion volume of the ith time interval respectively.

Figure 1. Schematic diagram of the experimental setup. 1: Intelligence control air compressor; 2: Pressure gauge; 3: Air tank; 4: Spring safety valve; 5: Gas pressure regulator; 6: Gas flow controller; 7: Model Ladle; 8: Drain; 9: Overflow receiver; 10: Constant liquid-level water tank; 11: Intake pipe; 12: Overflow pipe; 13: Digital camera.

Results and Discussion

The Behavior of Slag Entrapment

Figure 2. The powder entrapment in ladle at the gas flowrate of 5.52 Nl/min.

The powder entrapment in the ladle was taken by a camera as shown in Figure 2. It can be seen that, under a large gas flowrate, the bubbles rise in a gas-liquid plume that creates a raised area at the bath surface and a slag eye comes into being, and the upwelling flow from the bubble plume

209

turns horizontally at the bath surface and pushes the slag layer to the periphery of the ladle. When the gas flowate exceeds a critical value, the powder entrapment occurs.

Table III showed the effect of gas flowrates on the powder entrapment in ladle. It can be seen that the gas flowrate of 0.66 Nl/min is the critical gas flowrate for powder entrapment. With those gas flowrates smaller than 0.66 Nl/min, the powder entrapment will not occur. While it occurs with those gas flowrates larger than that, and the frequency of powder entrapment is higher and higher with the increase of gas flowrate. When the gas flowrate increases to 3.31 Nl/min, the powder entrapment occurs continuously. Considering the bad effect of slag entrapment on the steel quality, the gas flowrate should be smaller than 0.66 Nl/min to have a good control of inclusion in the ladle.

Table III. The effect of gas flowrates on the powder entrapment in ladle Gas flowrate (Nl-min"

') 0.22 0.44 0.55 0.66 0.88 1.33 1.77 2.21 3.31 4.42 5.52 6.63

Whether powder entrapment occurring

No No No Yes Yes Yes Yes Yes Yes Yes Yes Yes

Frequency of powder entrapment (-min"1)

0 0 0 2 7 19 22 37

continuously continuously continuously continuously

The Behavior of Inclusion Removal

Influence of Gas Blowing Time on Inclusion Removal. Inclusions removal rates at different time intervals under different gas flowrates were shown in Figure 3. It can be seen that most of inclusions can be removed within eight minutes and the inclusion removal rates at those intervals of eight to twelve, twelve to sixteen, sixteen to twenty, twenty to twenty-four and twenty-four to twenty-eight minutes decreased in turn. And the fastest removal rate occurred at the interval from zero to four minutes. Moreover, it can be found that the inclusion removal rate is less than 4.5% at the seventh interval, that is to say, almost all the inclusions which have the possibility to be removed disappear from the system within twenty-eight minutes blowing and then after the inclusion removal is not efficient any more.

Influence of Gas Flowrate on Inclusion Removal. Figure 4 shows the variation of inclusions removal rate with time under various gas flowrates. It can be seen that the gas flowrates of 0.22 Nl/min and 0.66 Nl/min show better efficiency on the inclusion removal than the other gas flowrates, and the inclusion removal rates under the other gas flowrates are all poor. Moreover, there is little difference on the inclusion removal under the other gas flowrates.

210

Figure 3. Inclusions removal rate at different Figure 4. The variation of inclusions removal time intervals under different gas flowrates. rate with time under various gas flowrates.

To get a better understanding on the effect of the gas flowrate on inclusion removal, Figure 5 shows the final inclusions removal rate within twenty-eight minutes vs gas flowrate. It can be seen that the inclusion removal rate decreased first and then increased with the increase of gas flowrate. And the gas flowrates of 0.22 Nl/min and 0.66 Nl/min show the best for inclusion removal in twenty-eight minutes blowing. But the gas flowrate of 0.66 Nl/min is the critical gas flowrate for powder entrapment. So, the gas flowrate of 0.22 Nl/min shows the best one for the inclusion removal since it will not cause powder entrapment.

Figure 5. The final inclusions removal rate within twenty-eight minutes vs gas flowrate.

Law on the Process of Inclusion Removal. Figure 6 shows the variation of inclusion concentrations with gas blowing time under two typical flowrates. It seems that the variation of inclusion concentrations with gas blowing time fits well with the exponent equation of C/C0=exp(M/) ,where M is a constant. That is to say, the inclusion removal was in exponential relationship with gas blowing time in a gas-stirred ladle.

211

Figure 6. The variation of inclusion concentrations with gas blowing time under typical gas flowrates of (a) 0.22 Nl/min and (b) 0.66 Nl/min (R is correlation coefficient).

Conclusions

(1) The gas flowrate of 0.66 Nl/min is the critical gas flowrate for powder entrapment. With those gas flowrates larger than that, the powder entrapment occurs and the frequency of powder entrapment is higher and higher with the increase of gas flowrate. When the gas flowrate increases to 3.31 Nl/min, the powder entrapment occurs continuously. (2) Most of the inclusions can be removed in eight minutes, and the fastest removal rate occurs at the interval from zero to four minutes. All the inclusions which have the possibility to be removed almost disappeared from the system in twenty-eight minutes. (3) The inclusion removal rate decreased first and then increased with the increase of gas flowrate. And the gas flowrate of 0.22 Nl/min shows the best one for the inclusion removal. (4) Inclusion removal was in exponential relationship with gas blowing time in gas-stirred ladle.

Acknowledgements

The present study was supported by the Fundamental Research Funds for the Central Universities (100402017) and National Outstanding Young Scientist Foundation of China (50925415).

References

1. S. Joo and R. I. L. Guthrie, "Modeling Flows and Mixing in Steelmaking Ladles Designed for Single- and Dual-Plug Bubbling Operations," Metallurgical Transactions B, 23 (6) (1992), 765-778. 2. S. Ganguly and S. Chakraborty, "Numerical Modelling Studies of Flow and Mixing Phenomena in Gas Stirred Steel Ladles," Ironmaking and Steelmaking, 35 (7) (2008), 524-530. 3. M. Y. Zhu et al., "Fluid-Flow and Mixing Phenomena in the Ladle Stirred by Argon through Multi-Tuyere," ISIJInternational, 35 (5) (1995), 472-479. 4. J. Mandai et al., "Mixing Time and Correlation for Ladles Stirred with Dual Porous Plugs," Metallurgical and Materials Transactions B, 36 (4) (2005), 479-487. 5. Z. R. Xu et al., "Investigation on Nonmetallic Inclusion of LACK Steel in LF Secondary Refining," Iron and Steel, 45 (10) (2010), 33-36.

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6. C. Fan et al., "Research on Cleanliness of Steel Grade 45 Produced by BOF-LF-CC Process," Iron and Steel, 38 (3) (2003), 18-20. 7. M. Söder, P. Jönsson, and L. Jonsson, "Inclusion Growth and Removal in Gas-Stirred Ladles," Steel Research International, 75 (2) (2004), 128-138. 8. Y. J. Kwon, J. Zhang, and H. G. Lee, "A CFD-Based Nucleation-Growth-Removal Model for Inclusion Behavior in a Gas-Agitated Ladle During Molten Steel Deoxidation," ISIJ International, 48 (7) (2008), 891-900. 9. M. Y. Zhu et al., "Numerical Simulation of Nonmetallic Inclusions in Gas-Stirred Ladles," Steel Research International, 76 (10) (2005), 718-722. 10. S. G. Zheng and M. Y. Zhu, "Water Model Study on Removing Inclusions in a Ladle with Argon Injected Through Nozzle and Porous Plug," Ada Metallurgica Sinica, 42 (11) (2006), 1143-1148. 11. S. G. Zheng and M. Y. Zhu, "Physical Modelling of Inclusion Behaviour in Secondary Refining with Argon Blowing," Steel Research International, 79 (9) (2008), 685-690. 12. S. G. Zheng, M. Y. Zhu, and S. S. Pan, "Experimental Study on the Behavior of Inclusions in RH Degasser," Ada Metallurgica Sinica, 42 (6) (2006), 657-661. 13. M. Y. Zhu and Z. Q. Xiao, Maths-physical Modeling of Steel Refining Process (Beijing, China: Metallurgical Industry Press, 1998), 124-125.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THE EFFECT OF HIGH SUPERHEAT ON THE SOLIDIFICATION STRUCTURE AND CARBON SEGREGATION OF FERRITE-BASED ALLOY

Honggang Zhong ', Yi Tan ', Huigai Li ', Xinping Mao 2, Qijie Zhai '

1 Shanghai Key Laboratory of Modern Metallurgy & Materials Processing. Shanghai University, Shanghai 200072, China

2 Guangzhou Zhujiang Iron and Steel Co., Ltd., Guangzhou 510730, China

Key words: High carbon steel, High superheat, Solidification structure, Macrosegregation

Abstract

The Ferrite-based alloy was melted and poured in-situ to investigate the effect of high superheat on the solidification structure and macrosegregation of permanent mold casting. The alloy solidified horizontally from the chill wall in a furnace, which be shut down after the alloy liquid be poured. The coarse equiaxed dendritic grains, which diameters are more than 5mm, are observed with high superheat (80-120 K), and the secondary dendritic arm spacing increases with superheat increasing. However, the degree of carbon segregation K(C), which is defined as the ratio of concentration of solute (C,) and average concentration of solute (Cme), are less than 1.2, which indicate the fluctuations of carbon concentration are small.

Introduction

High carbon steel tends to form a large mushy zone due to the wide solid/liquid temperature range which generally would result in serious porosity and macrosegregation defects [1, 2]. During solidification process of alloys, solute elements usually are rejected by primary dendrites, enriching in the interdendritic liquid. As the enriched liquid freezing, segregation occurred. Although macrosegregation could be reduced to some extent by following heat treatment, it is harmful to mechanical performance. High carbon steel is commonly produced by continuous casting and permanent mold casting because the rapid cooling rate is conducive to refine solidification structure and reduce the macrosegregation [3-5]. In addition, low superheat pouring is always suggested for the same purpose [1,6]. However, the low superheat steel liquid will lead to nozzle clogging or accidentally freeze in the tundish [7].

The solidification behavior is very important to the quality of slab and ingot. In continuous casting and permanent mold casting, the crystals usually grew in horizontal orientation due to the heat is transferred from side walls. However, crystal growth orientation is vertical, which now usually employed to physically simulate metal/alloy solidification [8, 9], During the growth of crystals in vertical direction, solutes often pile up on the horizontal solid/liquid (S/L) interface due to the gravity, and the heat transfer and liquid flow are different to horizontal solidifying

215

samples [10, 11]. To investigate the macrosegregation of solute, the horizontal dendritic growth direction is a better choice than vertical direction. In practice, the heat is mainly transferred in one dimension during solidification of slab or large-scale ingot, and the columnar grains grow nearly in a horizontal orientation. The improved horizontal unidirectional solidification technic could be used to simulate the solidification process of that. Recently, we have produced a horizontal dendritic growth setup to simulate the growth of dendritic by control the cooling water flux and furnace cooling rate.

Generally, it is believed that high pouring temperature is responsible for the coarse solidification structure and severe macrosegregation of carbon steel [1], whereas the high cooling rate could refine the grain structure and reduce the segregation. In this work, the effect of high superheat combined with water cooling on the solidification structure and solute macrosegregation in 1.8 mass% C steel were investigated by using the improved horizontal unidirectional solidification setup.

Experimental procedure

The high carbon steel used in the present study was a 0.65 mass% C steel rod. The sample weight was about 60g. To get 1.8 mass% C ferrite-based alloy, the rod was enclosed in an AI2O3 crucible with the same volume of carbonaceous casting powder, which could not only increase the carbon content but also prevent oxidization. The crucible was fixed on a water cooled pipe and was heated to 1873K by a MoSi2 furnace to melt the sample, and then cooled to pouring temperature. After holding for 30min, the melt was poured in-situ to contact the chill copper wall, which was cooled by running water in the pipe. After the melt was poured, the furnace was shut down and sample was solidified directionally against the chill wall. As the furnace temperature under 1273K, the furnace was pushed back and the sample was air cooled to room temperature. The solidified sample had a rectangular prism shape with a height of 20mm, a length of 40mm and a thickness of 10mm.

The solidified sample was then sectioned transversely at the middle of thickness and etched using saturation picric acid solution at 343K for about 30s. Metallographic images were observed using a Leica DM6000M optical microscope. The contents of C/S were measured by Leco CS600 carbon and sulfur analyzer.

The reasons to select the 1.8 mass% C ferrite-based alloy are as follows: 1) it has a wide solidification range (i. e. difference between the liquidus and solidus temperature), which is about 193K, as shown in Fig.l. The mushy zone of this steel is very deep during solidification. As the equilibrium partition ratio (K) of carbon in ferrite is about 0.25, it easily leads to severe segregation. If the segregation and solidification structure could be refined by high superheat, this method should be appropriate for other carbon steels. 2) The liquidus of this alloy is 1667K, which is much lower than that of the other carbon steels and the limit of MoSi2 furnace. It is tolerable for the furnace to acquire high superheat of these samples. Table I shows the mean

216

contents of carbon and sulfur of high superheat samples. The contents of carbon fluctuated from 1.65 to 2.00 mass%. We can still call this alloy as high carbon steel.

Fig. 1 Phase diagram of Fe-C alloy.

Table I. The Average Contents of Compositions in High Superheat Samples.

composition

Carbon

Sulfur

Average solute Content (mass %) A-80K

1.73

0.027

B-90K

2.00

0.047

C-100K

1.87

0.041

D-110K

1.824

0.034

E-120K

1.65

0.029

Results & Discussions

Fig.2 shows the solidification structure of high superheat samples. The columnar to equiaxed transition (CET) was investigated in samples A-C; while the super heat over 11 OK were transcrystalline structure. The secondary dendritic arm spacing (SDAS) were measured every 10mm from the chill zone (i. e. the surface of ingots) and drawn in Fig.3 (a). Due to the quick cooling rate against the chill wall, the SDASs of chill zone in each sample were less than that of columnar or equiaxed zone. And they increased with the superheat increasing from 80K tol20K. As the dendrites of chill zone in sample B & D could not be identified clearly, the data of their SDASs wasn't adopted.

The SDASs of sample E were much coarser than other samples; they were all greater than 100 urn. It suggested that the running water (2L/min) in chill wall was not sufficient for the 120K superheat melt. The chill zone was heated by sensible and latent heat; and the dendrites got

217

coarse due to this heat treating. Due to the coarse SDAS, high order dendrites developed in the interstice of secondary dendrites (see the coarse columnar dendrites in Fig.2 (c)). In fact, as the high order dendrites filled the interstices, the solidification structure of sample E was as fine as sample A and few shrinkage voids emerged in coarse columnar zone.

Fig.2 Dendritic structure of high carbon steel, (a) A-80K, (b) C-100K, (c) E-120K.

Local solidification time (tj) or mean cooing rate (s) during solidification is related to SDAS (fa) [10],

Z1=a{t/y=b£- (1)

Where a, b and n are constant. Jacobi [12] acquired b= 109.2 and n=0.44 by steady state directional freezing experiments of high carbon steel,

^ = 109.2(e)"044 (2)

The mean cooling rates e=1.04K/s and £=5.98 K/s can be calculated assuming SDASs to be 107.3 \im (chill zone of sample E) and 49.7 um (chill zone of sample A) separately. It suggests the high super heat restrained the heat transfer and delayed the solidification of steel in chill zone. The mean cooling rates e^.OS-O.SlK/s when SDAS are 80-120 um in hot zone of samples. It suggests the cooling rate of sample reduces with superheat dramatically. The calculated solidification time of sample E is twice longer than sample A.

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Fig.3 (a) Variation of SDASs from chill wall to center (b) The rate of volume with porosity to sample volume.

The rate of volume with porosity (see Fig.2 (b)) to sample volume versus the superheat is shown in Fig.3 (b). This trend is not the same as that of SDAS with superheat, which generally increases with superheat increasing, but shows a peak at 100K superheat. The shrinkage porosity distributed over 20% volume of sample. Dense ingots could be obtained when superheat is 80K or 120K. in these tests.

Fig.4 Carbon macrosegregation of high superheat samples.

The carbon macrosegregation is expressed as a ratio Kc, which is the ratio of carbon content C, of the specimens to the average value C™, of the sample. Positive and negative segregation refer compositions above and below the average composition, respectively:

219

Kc=-pr~ (3)

C^=-tq (4) « w

The degree of macrosegregation (ATc) was presented in Fig.4. In these tests, all Kc ranges from 0.79 to 1.16 along the whole samples. It means the segregation wasn't severe. However, both positive and negative carbon segregation were investigated in chill zone (surface of sample). This phenomenon has not got a clear explanation.

The carbon segregations of sample A (superheat=80K) was increased linearly with position increasing, and the fluctuant range over 0.4. As the macrosegregation could hardly be reduced by the following heat treating and rolling process, low superheat was not favorable to production in these tests.

Conclusions

The effect of high super heat (80-120K) on the solidification structure and solute macrosegregation in high carbon steel were investigated by horizontal unidirectional solidification technic. The melt was poured in-situ and solidified against the chill wall, which was cooled by running water. The main results of this investigation are summarized as follows:

1) The carbon segregation reduces with superheat increasing, whereas SDAS increases, which reflecting local cooling rate and solidification time. The solidification time increases with superheat increasing significantly.

2) The porosity rate, which is defined as the ratio of volume with porosity to sample volume, is over 20% when superheat is 100K. However, tense ingots with few voids are obtained when the superheat are 80K. and 120K.

3) The secondary dendritic arms of sample poured with superheat 120K were much coarser than lower superheat samples, but the high order dendrites developed in the interstices and produced fine solidification structure.

4) Normal carbon segregations are observed in sample A (superheat=80K), and the fluctuant range over 0.4.

In conclusion, High superheat pouring method combined with water cooling could reduce solute segregation and acquire uniform solidification structure of high carbon steel. For these tests, the optimized superheat for 1.8 mass% C steel is 120K.

Acknowledgments

The authors gratefully acknowledge the financial supports from the National Natural Science Foundation of China (Grant Nos. 50734008 and 50904044).

220

References

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

REFINEMENT OF LIGAMENTS OF NANOPOROUS Ag RIBBONS

BY CONTROLLING THE SURFACE DIFFUSION OF Ag

Tingting Song1,2, Yulai Gao1'2*, ZhonghuaZhang31, Qijie Zhai2

'Laboratory for Microstructures, Shanghai University, 99 Shangda Road; Shanghai 200436, P. R. China

2School of Materials Science and Engineering, Shanghai University, 149 Yanchang Road; Shanghai 200072, P. R. China

3Key Laboratory for Liquid-Solid Structural Evolution and Processing of Materials (Ministry of Education),

School of Materials Science and Engineering, Shandong University, 17923 Jingshi Road; Jinan 250061, P. R. China

Keywords: Nanoporous silver, Chemical dealloying, Finer ligament, Surfactant

Abstract

Nanoporous silver (NPS) with different nanoporosity was manufactured by chemical dealloying of rapid solidified Al-Ag ribbons consisting of two distinct phases of a-Al(Ag) and Ag2Al. The as-dealloyed samples were characterized by X-ray diffraction (XRD), scanning electron microscopy (SEM), and energy dispersive X-ray (EDX) analysis. It is found that the width of the average ligaments can be dramatically decreased and more homogeneous when surfactants were added to the H2SO4 solution, which would definitely increase the specific area of NPS and this is pretty desirable for certain applications, such as catalyst, absorption, etc. The finer ligaments were attributing to the hindered surface diffusion of Ag atoms by surfactants, and the surface diffusion coefficient (Ds) of Ag was decreased from the order of 10"14 to 10"16 m2/s as a result of surfactants. This simple method to reduce the ligament width of NPS is also anticipated to prepare other related nanoporous metals.

Introduction

Nanoporous metals have attracted increasingly interest in recent years contributing to their wide potential applications in catalysts,[l] sensors,[2] actuators,[3] and so forth. Nanoporous metals are commonly fabricated by template methods[4] which are generally technically difficult, time consuming and template-sacrificed. Recently, dealloying [5] as

Corresponding author. Tel./fax: +86-21-56332144. Email: ylgao(_!shu.edu.cn. f Conesponding author. Tel./fax: +86-531-88396978. Email: z.h_zhang@sdu.edu.cn.

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a simple and effective strategy is widely used to produce nanoporous metals with random porous structure. The so-called dealloying is a corrosion process during which the less noble (LN) component is dissolved away from the precursor alloy, and simultaneously the remained more noble (MN) element diffuses and agglomerates into a well-defined 3D bi-continuous nanoporous structure.[5] To date, except for the prototypical Au-Ag [5] system, the dealloying process has been studied in other various alloy systems such as Ni-Cu,[6] Cu-Au,[7] Pt-Si,[8] and correspondingly can be utilized to prepare nanoporous metals and alloys.

It is well known that silver is wildly used as excellent catalysts[9] and substrate for surface enhanced Raman scattering. [10] Nanoporous silver (NPS) with random porous structure and ultrahigh surface area would definitely make silver more useful in such fields. Jia et al.[11] and Yeh et al.[12] have prepared NPS through an electrochemical alloying/dealloying process. In our previous work, NPS have been fabricated by dealloying of Al-Ag alloys in various etching solutions. [13] For applications that require ultrahigh surface areas, a coarsened ligament structure may be undesirable and it is urgent to minimize the ligament size of nanoporous metals including NPS.[14] It is accepted that the formation and resultant ligament size of the nanoporous metals are mainly determined by the diffusion of NM component, which can be expressed by the diffusion coefficient (A,). [15] So controlling the diffusion of NM component is anticipated to minimize the ligament size. Many methods have been developed to control the ligament size of nanoporous metals, such as low temperature dealloying[16] to control the applied potential of electrochemical dealloying[17] and introduction of additives in the alloy[15] by adding surfactants into the etching solution[18] to slow down the atomic diffusion. Among them, low temperature dealloying and potential-modified dealloying are technically difficult and complicated for practical application. In addition, adding a third element in the alloys is economic costly as the third element which would restrain the diffusion of MN element in the alloy. Surfactants are generally used as stabilizer for preparing metallic nanoparticles. Silver (8-32 nm) and gold (6-21 nm) nanoparticles have been synthesized by Hiramatsu et al[19] using oleylamine as stabilizer. In preparing Pd- and Pt- nanoparticles, Chen et al[20] have also employed oleic acid and oleylamine as stabilizers to control the surface diffusion of Pd and Pt atoms.

In the present study, Al-Ag alloys are chosen as precursors, H2S04 solution as etching solution and oleic acid and oleylamine as stabilizers to investigate the dealloying behavior of Al-Ag alloys with and without surfactants controlling the diffusion of Ag.

Experimental section

The alloy with nominal compositions of Al-25(at.%)Ag were prepared from Al (purity, 99.99 wt.%) and Ag (purity, 99.99 wt.%) being rapidly solidified into ribbons with a cross section of (0.025 ~ 0.045) x 4 mm2 using a single roller melt-spinning method.

224

Dealloying of these rapidly solidified Al-Ag alloy ribbons was carried out in a 25 vol.% H2SO4 solution (100 ml) and 25 vol.% H2S04 solution (100 ml) with oleic acid (C17H33COOH, 90% purity, Alfa Aesar, USA) and oleylamine (Ci8H35NH2, 70% pure, Alfa Aesar, USA) (2 ml each), respectively. And dealloying was implemented at 371 ± 2 K until no visible bubbles escaped. The as-dealloyed samples were then removed from the acid solution and rinsed in distilled water and dehydrated alcohol several times respectively to guarantee their cleanness.

X-ray diffraction (XRD) patterns of the rapidly solidified Al-Ag alloys and the as-dealloyed samples were collected using a Rigaku D/Max-2200 X-ray diffractometer with Cu Ka radiation (k = 0.154056 nm). The microstructures and composition of the NPS samples were observed by scanning electron microscopy (SEM: JSM-6700F by Hitachi High-Technologies Corp.) and an energy dispersive X-ray spectroscopy (EDX: INCA by Oxford).

Results and discussion

Figure 1 shows the XRD patterns of the rapidly solidified Al-25Ag alloy ribbons and the as-dealloyed Al-25Ag samples from two solutions (100 ml 25 vol.% H2SO4, 100 ml H2SO4 + 2 ml C17H33COOH + 2 ml C18H35NH2). It is clear that two phases exist in the starting Al-25Ag alloys: a solid solution a-Al(Ag) and intermetallic compound Ag2Al, and that a-Al(Ag) phase is dominant by comparing intensities of diffraction peaks of both phases.[13] Yamauchi et al[21] and Wang et al[22] have proved the important role of an a-Al(Ag) phase during the dealloying of bi-phase Al-Ag alloys and the formation NPS as well. In addition, it can be seen from Figure 1 that after dealloying in two solutions, only a single Ag phase can be identified in the as-dealloyed samples, by comparing all the diffraction peaks with standard lines of Ag (PDF file, # 04-0783). These results indicate that rapid solidified bi-phase Al-Ag alloys can be completely dealloyed in H2S04

solution with and without surfactant. It should be noted that the dealloying process lasts for different duration in different solutions (see Table 1) and dealloying in H2S04 with and without surfactant takes 3 h and 1 h, respectively. As the oleate ions (C17H33COO) can easily adsorb on the surface of noble metals,[18] the reaction between Al and H2S04

would be prevented to some extent. In addition, when the acid concentration turns from 5wt.% to 25 vol.%, the dealloying time for Al-25Ag ribbons changes from 5 h in our previous study[13] to 1 h here which clarify that the acid concentration is definitely vital for the dealloying process.

Figure 2 and Figure 3 shows the microstructure of the as-dealloyed Al-25Ag ribbons. From the plan view of NPS from H2S04 (see Figure 2(a, b)) we can see a homogeneous nanoporous structure without clear cracks over a large area. The cross section view in Figure 2(c) shows well defined structure with excellent mechanical integrality which indicates the whole thickness has gone through the same dealloying process. The high magnified cross section view (see Figure 2(d)) display the typical 3D bi-continuous

225

nanoporous microstructure with an average ligament width of 284 ± 74 nm (listed in Tablel). In this article, at least 30 ligaments were measured from the typical SEM images

Figure 1. XRD patterns of Al-25 Ag precursors and the as-dealloyed samples in H2SO4 solutions with and without surfactants (C17H33COOH + C18H35NH2).

in the second electron mode (SEI images) to get a statistical determination of their size. In addition, the typical EDS result is shown in Figure 2(e) and only Ag traces can be observed corresponding well to the XRD result in Figure 1.

Table 1. The dealloying conditions of Al-25 at. %Ag ribbons and the ligament width and Ds of the resultant NPS.

Alloy

Al-25 at%Ag

Al-25 at.%Ag

Temperature (K)

371 ±2

371 ±2

Dealloying time (h)

1

3

Surfactant

No

Yes

Ligament width (nm)

284 ± 74

159 ±36

D, (m2/s)

1014

io-1 6

Figure 3(a, b) are the plan views of NPS dealloyed for 3 h in H2SO4 with surfactants, from which the surface with homogeneous microstructure and good mechanical integrity can be observed. The cross section view in Figure 3(c) shows well refined structure without clear cracks either. The high magnified cross section view (see Figure 3(d)) display the typical 3D bi-continuous nanoporous microstructure with an average ligament width of 159 ± 36 nm (see Table 1) much finer than the 284 ± 74 nm above. The typical EDS result in Figure 3(e) with only Ag traces clarifies the completeness of dealloying further.

The dealloying conditions for the two dealloying processes are shown in table 1. It is interesting that the average ligaments size of NPS decreases dramatically from 284 ± 74 nm to 159 ± 36 nm as surfactant was added to the acid solution. It is seemed that the surface diffusion of Ag atoms, which determines the formation of porous structure,[16] has been hindered when the C17H33COOH and C18H35NH2 was added into the H2SO4

226

etching solution. The surfactant is frequently employed in the production of nanoparticles,[19, 20,23] which is usually considered as the stabilizer by preventing the

Figure 2. Microstructure ((a, b): plan view, (c, d): cross-section view) and typical EDS result (e) of NPS dealloyed from 25 vol.% H2S04 solution.

surface diffusion of the nanoparticles. The molecules in surfactant can attach to the metal surfaces via physisorption due to vander Waals attraction[24] or via chemisorption with covalent bonding.[25] And then the surface diffusion of metal adatoms would be hindered by molecules or ions attached to the Ag surfaces. Obviously, the surface diffusion of Ag atoms in the two processes should finally be responsible for the difference in the size of the resultant NPS. Based on the surface

227

diffusion controlled coarsening mechanism, Ds can be estimated by the following equation: [26, 27]

Figure 3. Microstructure ((a, b): plan view, (c, d): cross-section view) and typical EDS result (e) of NPS dealloyed from 25 vol.% H2SO4 solution with surfactants.

D =[d(t)TkT (1) 32yta

where k is the Boltzmann constant (1.3086 x 10"23 J K"1), y is the surface energy (1 J m"2),[28] d(t) is the ligament size of the as-dealloyed Al-15Ag alloy at the dealloying time t, a is the lattice parameter of Ag (4.086 * 10"'° m V ) and T is the dealloying temperature (363 K). According to the parameters in Table 1, the Ds values of Ag atoms

228

were calculated for the dealloying of the Al-Ag alloys in H2SO4 solutions with and without surfactant, and are listed in Table 1. It is clear that Ds of Ag atoms in etching solution with surfactant (~1016 m V ) is 2 orders less than that only in etching solution (~10"14 m V ) which directly clarify that the adding of surfactant restrain the surface diffusion of Ag atoms and make the ligaments of nanoporous structure much finer and more homogeneous.

Conclusions

Monolithic nanoporous silver can be manufactured by dealloying of Al-25Ag ribbons in 25 vol.% H2SO4 solution. The average ligament size of the resultant NPS can be dramatically decreased from 284 nm to 159 nm by adding surfactant to the acid solution, which is beneficial for the increase of the specific surface area. It is proposed that the adding of surfactant into the etching solution would restrain the surface diffusion of Ag atoms and the calculated Ds (~10'16 m V and ~10"14 m V ) directly proved it. This simple procedure for minimizing the average ligament size and increasing the specific surface area is desired for the wide potential applications of nanoporous metals and can be applied in other alloy systems.

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Nanoparticles with Adjustable Sizes and with Exchangeable Surfactants," Chemistry of Materials, 16 (13) (2004),

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THERMAL ANALYSIS OF THE COMPOSITION OF POLY(ACRYLIC ACID)/CARBOXYMETHYLSTARCH USED AS A POLYMERIC BINDER

Beata Grabowska1*, Mariusz Holtzer1, Sonja Eichholz2, Krzysztof Hodor2, Ewa Olejnik'

'AGH - University of Science and Technology, Faculty of Foundry Engineering, Reymonta 23, 30 059 Krakow, Poland,

applications Laboratory Thermal Analysis, NETZSCH-Gerätebau GmbH, Wittelsbacherstr. 42, 95100 Selb/Bavaria, Germany

Keywords: thermal degradation, poly(acrylic acid), carboxymethylstarch, polymer composition, polymeric binder

Abstract

Samples of poly(acrylic acid)/carboxymethylstarch used as a binding agent in molding sands were investigated. Methods of thermal analysis (DTG, TG, DSC) were applied to assess the thermal stability of the investigated polymer sample by estimation of temperature and thermal effects of transformations occurring during its heating. In the temperature range -100-1000°C none polymorphic transformations were found. It was established that the degradation process starts at a temperature app. of 130°C. On the bases of the analysis of volatile products of the polymer decomposition performed by means of the IR spectral method and the thermo gravimetric (TG) method coupled 'on-line' with the mass spectrometry (MS) the signals for low molecular masses were found in the temperature range: 300-400°C. This indicates that the degradation process occurs, polymer chains are undergoing fragmentation and low molecular compounds are formed. In the final stage H2O and CO2 are formed. At the higher temperature range (350-600°C) compounds and alkyl radicals of higher mass numbers are formed.

Introduction

All materials, raw materials and products, which in a direct or indirect way are used for making and finishing casting moulds, are considered moulding materials [1]. They can be divided into materials being matrices of moulding sands and protective coatings and materials binding matrix grains (binding agents). Binding agents provide a matrix with a determined mechanical strength, allowing to obtain castings of the required shapes and dimensions. On account of their chemical character binding agents are divided into inorganic (among others: clays, cements, water-glass and bentonites) and organic (among others: phenol-formaldehyde resins, carbamide resins) binders. However, moulding sands with inorganic binders have several significant drawbacks, such as: long drying and hardening time, hygroscopicity, cracking of ready moulds, difficult knocking out, weak reclaimability and low strength. To this end, their applicability is limited mainly to large castings of simple shapes. On account of inorganic binders faults, mentioned above, more and more often moulding sands based on organic binders - including polymers - are applied. These sands are characterised by good technological properties such as: proper strength, good fluidity and working time, which enable obtaining castings of the required functional parameters. Due to the fact, that some synthetic polymers constitute environmental hazards, biopolymers - being reclaimed natural polymers - seem interesting [1-4].

231

Investigations on an application of polymers, including biopolymers, as ecological binding agents have been carried out for some years already in the Laboratory of Environmental Protection of the Department of Foundry Engineering, AGH [5-7]. So far the polyacrylic binder and a way of its hardening in a moulding sand was worked out as an ecological alternative of other binding materials [8], Further investigations are carried out with the aim of finding new binders or modifying known ones to limit their harmfulness for the environment and to improve their moulding sand technological properties. In order to attain this purpose it is necessary to know, not only physical and chemical properties - including toxicity - of the initial material (polymer binder), but also mechanisms of its thermal degradation [9, 10]. Spectroscopic FTIR methods and thermal analysis (DSC, DTG, TG) were applied in this study to determine the thermal stability of a polymer binder in a form of a polymer composition: poly(acrylic acid)/carboxymethylstarch, by finding the destruction temperature and thermal effects of transformations occurring during heating. In addition, the analyses of volatile decomposition products, by means of the thermogravimetric (TG) method coupled on-line with the mass spectrometry (MS), were performed. All these investigations allowed to explain the polymer binder degradation process, which is of an essential meaning in respect of its utilisation as the moulding sand binding agent, with taking into consideration its contact conditions with liquid metal (temperature app. 1400°C).

1. Research methods Materials Polymer binder (Komp A): a polymer composition made up of a synthetic polymer: poly(acrylic acid) produced by BASF (Fig. 1) and a natural polymer: carboxymethylstarch by Xenon®, (Fig. 2) in the 1:1 weight ratio, formed by mixing in water.

c=o OH

Figure 1. Poly(acrylic acid) structure.

CH2-OH ÇH,-0-CH2-COOH

-O

Figure 2. Carboxymethylstarch structure.

Preparing samples for examinations Cross-linking and hardening samples with microwaves: RM 2001 Pc microwave reactor by Plazmatronika (800 W, 90 s).

232

Infrared spectroscopic examinations. FTIR Infrared spectroscopic examinations were carried out using a Digilab Excalibur FTS 3000 Mx spectrometer with a DTGS detector, electrically cooled. The spectrometer is equipped with two attachments: ATR with a ZnSe crystal for multiple reflections and a transmission attachment.

Thermal analysis examinations The NETZSCH model STA 449 F3 Jupiter® simultaneous thermal analyzer can be used to measure the mass change and transformation energetics of a wide range of materials. The top-loading STA can be equipped with various easy exchangeable TG, TG-DTA or true TG-DSC sensors and with different furnaces to accommodate different application areas. The system employed for this work was equipped with a steel furnace capable of operation from -150 to 1000°C. For control of the measurements as well as for data acquisition, modern digital electronics and the well established NETZSCH PROTEUS' 32-bit Software are employed. Several Advanced Software packages like c-DTA' (calculated DTA-signal), Super-Res* (rate-controlled mass change) or Thermokinetics- are available. Furthermore, combining both thermogravimetric and spectroscopic methods such as MS enables identification of the evolved gases.

Microscopic examinations Microscopic examinations were performed by means of the optical microscope of the Nikon Company, ECLIPSE LV 100 with objectives: Nikon Lu Plan Fluor 20x/0.45A and Nikon Lu Plan ELWD 50x/0.55B.

2. FTIR examinations and their analysis

In order to elucidate processes occurring in the examined polymer composition poly(acrylic acid)/carboxymethylstarch under an influence of heating the spectroscopic FTIR examinations were performed. Temperature spectra were made for the polymer composition sample in the temperature range: 25-200°C (operation range of die temperature attachment of the IR spectrometer) and in the temperature range: 250-500 °C (heating furnace FCF ZZHM of the CZYLOK Company). The heating process (25-200°C) was performed in a continuous way, and the spectra were recorded at the required temperatures. In the temperature range: 25-300°C in the FTIR spectra (Fig. 3), within wave numbers 3700-3000 cm"1, the wide band - corresponding to stretching vibrations of the hydroxyl group (band of a free -OH group from water and hydrogen bonds: -O-H 0=C-) - occurs.

233

Figure 3. FT1R temperature spectra for the polymer composition poly(acrylic acid)/carboxymethylstarch.

Gradually with the temperature increase the intensity of this band decreases, which is related to the solvent water evaporation and constitutional water release during heating. At a temperature of 100°C the band corresponding to stretching vibrations of-C-H occurs near 2950 cm"1. As a

temperature increases this band shifts in the direction of lower wave numbers -*v=2930 cm"1, and its intensity increases. In addition, at a temperature of 300°C it undergoes splitting. This fact can indicate the formation of new groups (organic compounds of short chains) with the participation of -C-H group. In the range of wave numbers 1700-1600cm' two bands corresponding successively to vibrations of the carbonyl -C=0 group and asymmetric carboxyl -COO'group are seen. The band characteristic for carboxyl groups gradually decays and at a temperature of 150°C only one band corresponding to vibrations of carbonyl (-C=0) group - of a significantly increased intensity - remains. This change is related to the progressing degradation process in polymer chains. Two bands present in the range: 1400-1450 cm"1, which correspond to deformation vibrations of -(CH2),, groups (symmetric shearing vibrations) shift in the direction of smaller wave numbers, when the temperature increases. At a temperature of 300°C the appearance of new bands, being the result of the progressing degradation process, can be noticed

in the range 2000-1000 cm"1. In addition, the newly formed band in the range v=2300 cm1 is related to the evolution of CO2. Finally, at a temperature of 500°C the majority of characteristic bands decays, thus it can be stated that practically the total destruction of the polymer binder occurred.

3. Thermal analysis examinations

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To explain more broadly the degradation process of the investigated polymer binding agent in a form of the polymer composition: poly(acrylic acid)/carboxymethylstarch the thermal analysis TG-DSC was performed. Figure 4 depicts the TG-DSC results for the sample polymer composition. At sub-ambient temperatures, no effects were observed. Three mass loss steps of about 7%, 27% and 41% occurred which were accompanied by endothermic effects visible in the DSC signal. Maxima in the rate of mass change occurred at about 140°C, 270°C and 420°C.

Figure 4. Temperature-dependent mass change (TG), rate of mass change (DTG, dashed line) and heat flow rate (DSC) of the sample polymer composition in the hardened state.

At the -100°C-0°C temperature range no polymorphic transformations were detected. The shape of thermal curves is complex because the degradation process progresses in stages, which is due to the structure and the physical chemistry properties of the two-component biopolymer composition in question. Shapes of thermal curves support the claim that the degradation process starts at the temperature of aboutl30°C. To analyse the volatile products of decomposition, the thermogravimetry (TG) method coupled online with mass spectrometry (MS) was used in this research (Fig. 5). Within the 150°C-400°C temperature range, signals of low molecular weights were detected, proving the occurrence of processes of degradation, polymer chain fragmentation and the formation of low molecular compounds (molar masses: 15, 18, 22, 37, 38, 39, 41, 43, 44, 45, 46, 51 and 55 g/mol), during which mainly the following low molecular compounds are formed: H2O, CO2 and alkyl radicles among others CH3' [10].

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Figure 5. Mass changes (TG) and MS signal of composition poly(acrylic acid)/carboxymethylstarch.

A microscopic photograph taken of a moulding sand sample with a polymer binder baked at the temperature of 500°C shows carbonisations of the binder resulting from the action of high temperature, which means that the polymer has suffered partial destruction (Fig. 6).

Figure 6. The morphology of a moulding sand with a polymer binder made using quartz sand (baking temperature of 500°C).

4. Degradation mechanism of the polymer binder

On the basis of the obtained results (FTIR, DSC-TG, MS) and their analysis, it can be stated that the thermal degradation of the tested binder occurs in two steps, in accordance with the general radical decomposition mechanism [11, 12]. This degradation is preceded by the dehydration process, which occurs at a temperature of app. 130°C and means loosing the solvent water first and then the constitutional one. The dehydration process can be accompanied by a formation of intermolecular hexamember anhydrite rings. The real thermal degradation process starts at the initiation stage, during which free macroradicals are formed as a result of breaking interatomic

236

bond in the main chain of the macromolecule and a formation of two active macroradicals at the ends of shorter chains, which is presented by the reaction below (1):

R i

R i

—CH2-CH-CH2-CH-

In addition, the following reactions can occur (2 -3):

-*- — ChL-CH + -CH.-CH 2 i

R (1)

-CH--CH— 2 i

c=o I

OR

-CH--CH— 2 i

c=o I

o

R'

(2)

-CH.-CH— 2 i

c=o I

OR

-*- CO + *OR

C02 + R* (3)

Where: R is a substituent.

Macroradicals formed in such a way can undergo the following reactions [11, 12]: ♦ Depolymerisation process (releasing monomer under the influence of a high

temperature), ♦ Transferring of the free radical, both in the intramolecular and extramolecular reactions, ♦ Disproportionate» reaction, during which the new macroradical and macromolecule

containing dual bond is formed, ♦ Mutual recombinations forming either branched or cross-linked structures.

Above a temperature of 300°C an intensive cracking of chemical bonds in macromolecules occurs, leading to tearing the majority of bonds. The process is very rapid with a monomer evolving, followed by volatile organic compounds, which finally leads to the sample destruction. The destruction products are gaseous substances mainly small-molecule aliphatic hydrocarbons such as e.g. methane, and CO2.

5. Summary

Thus, the spectroscopic infrared (FTIR) examination performed at increased temperatures confirmed the thermal analysis results (TG-DSC) as well as enabled to determine the thermal stability of polymer binding agent in a form of a composition: poly(acrylic acid)/ carboxymethylstarch. In addition, the performed thermal analysis can be related to a polymer binder behaviour in the technological process of producing castings. Essential information concerning the polymer binder degradation when in contact with liquid metal (binder destruction in moulding sands), as well as the destruction products generation

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(emissivity of non-toxic gaseous products into the atmosphere) constitute the knowledge supplement concerning the application of poly(acrylic acid)/carboxymethylstarch composition as binding agent in foundry technology.

Preparation of the paper has been supported by scientific grant under project N N 507326 836 (2009-2011).

References 1. Mariusz Holtzer, "Trends in sand moulds and cores with organic binder", Archives of Foundry Engineering, 9 (2003), 189-195. 2. Xia Zhou, Jinzong Yang, Guohui Qu, "Study on synthesis and properties of modified starch binder for foundry", Journal of Materials Processing Technology, 183 (2007), 407-412. 3. Mitch Patterson, Jerry Thiel, "Developing Bio-Urethanes for No-Bake", Foundry Managemente: Technology, 6 (2010), 14-17. 4. Manas Chanda, Salil K. Roy, Industrial Polymers, Specialty Polymers and Their Applications, (CRC Press, Taylor&Francis Group, 2008). 5. Beata Grabowska, Mariusz Holtzer, "Application of microwave radiation for crosslinking of sodium polyacrylate/silica gel system used as a binder in foundry sands", Polimery, 11/12 (2007), 841-847. 6. Beata Grabowska, Mariusz Holtzer, "Application of Spectroscopic Methods for Investigation of Cross-Linking Process of Sodium Polyacrylate by Various Methods", Polimery, 7/8 (2008), 531-537. 7. Beata Grabowska, Mariusz Holtzer, "Structural examination of the cross-linking reaction mechanism of polyacrylate binding agents", Archives of Metallurgy and Materials, 2 (2009), 427^137. 8. Beata Grabowska, Mariusz Holtzer, Patent PL207459B1 (2010). 9. Beata Grabowska, "Thermal degradation of biopolymer binders: the example of starch-poly(acrylic acid)", Archives of Foundry Engineering, 10 (2010), 221-224. 10. B. Grabowska et al., "Thermal analysis of new polymer binding agents with a biopolymer fraction, based on example of the BioCol binder", Archives of Metallurgy and Materials (in print 2011). 11. Sheng-Cong Liufti, Han-Ning Xiao, Yu-Ping Li, "Thermal analysis and degradation mechanism of polyacrylate/ZnO nanocomposites", Polymer Degradation and Stability, 87 (2005), 103-110. 12. Abdul M. Kader, Anil K. Bhowmick, "Thermal ageing, degradation and swelling of acrylate rubber, fluororubber and their blends containing polyfunctional acrylates", Polymer Degradation and Stability, 79 (2003), 283-295.

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Mechanical Behavior Related to Interface Physics

The proceedings contained in this section have not been edited or reviewed by the conference program organizers. The opinions and statements expressed in the proceedings are those of the authors only and are not necessarily those of the editors or TMS staff. No continuations or endorsements are intended or implied.

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INTERFACIAL STRENGTH OF Al/Zr/DU-10%wtMo SUBJECT TO DIFFERENT LOADING MODES

C. Liu, ML. Lovato, and W.R. Blumenthal

Materials Science and Technology Division Los Alamos National Laboratory, Los Alamos, NM 87545, USA

Keywords: Interfacial strength, Al/Zr/DU-10wt%Mo, Mixed-mode loading, D1C

Abstract

Compact tension (CT) experiments were conducted with fixtures that allowed mode-I (tensile opening mode), mode-II (shearing mode), and mixed-mode loading to measure the interfacial strength between HIP-clad Al and Al, and Al and Zr/DU-10wt%Mo. Specimens were made with the same HIP process used for making thin composite foils, but instead used 25 mm thick Al-6061 cladding that allowed specimens to be gripped without adhesives. Three configurations of specimens were tested: (1) Al/Al specimens with a pre-crack along the seam; (2) specimens containing both a Zr/DU-10wt%Mo layer and an Al/Al seam along part of the interface; and (3) specimens containing only a Zr/DU-10wt%Mo layer at the interface, but with a pre-notch along part of the interface. Digital image correlation (D1C) was used to measure full-field deformations during the test. The results show that mode-I loaded interfaces exhibit the weakest strength and the widest scatter. The strength increases when more shearing component is introduced.

Introduction

A series of compact tension (CT) experiments, combined with the Arcan loading fixture, were conducted to study and measure the interfacial strength between HIP-clad Al and Al, and Al and Zr/DU-10wt%Mo. The Arcan fixture allows different loading modes to be applied to a specimen from mode-I (tensile opening mode) to mode-II (shearing mode). Specimens were obtained by sandwiching a thin foil of Zr/DU-10wt%Mo between two thick (25mm) Al-6061 blocks and applying the same HIP process that is used for making thin (1mm) Al/Zr/DU-10wt%Mo foils. Three types of CT specimens were tested: (1) specimens with a pre-crack along an Al/Al only seam; (2) specimens containing both a Zr/DU-10wt%Mo layer and an Al/Al seam along part of the interface; and (3) specimens containing only a Zr/DU-10wt%Mo layer at the interface, but with a pre-notch along part of the interface. Three loading angles were tested: 0° (mode-I), 45° (mixed-mode), and 90° (mode-II shear). For the cases of 45° (mixed-mode) and 90° (mode-II), the specimens with Zr/DU-10wt%Mo layer in the middle and a pre-notch along one of the interfaces are also subject to either positive or negative shear. Digital image correlation (DIC) was used to obtain full-field displacements and strains during the test.

Experimental Set-up

The interfacial fracture toughness has been shown to strongly depend on the mixity at the crack tip [1]. The crack tip mode mixity is the ratio of the shearing deformation parallel to the interface to the opening deformation perpendicular to the interface.

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To investigate the effect of crack tip mode mixity on interfacial failure, we use the compact tension (CT) specimen combined with the Arcan loading fixture, as shown in Figure 1. The CT specimen, shown in Figure 1(a), has a rectangular shape and the Zr/DU-10wt%Mo layer is located at the center. A pre-crack is machined along one of the interfaces. The Arcan loading fixture has two identical pieces and each piece has a series of holes, where loading pins can be inserted. As shown in Figure 1(a), these holes form a circle and the pair of holes used for loading pins is a diameter of the circle. The initial interfacial crack tip is located at the center of the circle. By varying the pairing of pinholes, different mode mixities at the initial crack tip can be achieved, from pure opening to pure shearing. Meanwhile, by flipping the CT specimen, or equivalently by changing the orientation of the crack (toward right or toward left), both positive and negative shear can be achieved. Figure 1(b) presents the photograph of the actual setup, where the CT specimen is loaded in pure mode-II (shearing mode).

Figure 1 : (a) Schematic of the compact tension (CT) specimen and the Arcan fixture; (b) The setup.

Three different CT sample configurations are considered and these three sample types are shown in Figure 2: (a) Specimens with a pre-crack along the Al/Al seam (named as Al-Al-notch sample); (b) specimens containing both a Zr/DU-10wt%Mo layer and an Al/Al seam along part of the interface (named as Al-DU-corner sample); and (c) specimens containing only a Zr/DU-10wt%Mo layer at the interface, but with a pre-notch along part of the interface (named as Al-DU-notch sample). The specimens shown in Figure 2(a) are used for measuring the strength of

Figure 2: Three types of CT samples tested: (a) Al-Al-notch (b) Al-DU-corner and (c) Al-DU-notch.

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the bonding of Al and Al from the HIPing process, the samples shown in Figure 2(b) are for studying the strength of material at the tip of the Zr/DU-10wt%Mo layer, and the samples shown in Figure 2(c) are for measuring the strength and toughness of the Al/Zr/DU- 10wt%Mo interface, where an interfacial crack is present. Three different loading angles are also considered: 0° (mode-I), 45° (mixed-mode), and 90° (mode-II shear). For the cases of 45° (mixed-mode) and 90° (mode-II), the specimens shown in Figure 2(c) can be subject to either positive or negative shear depending on the location or the orientation of the pre-notch.

In this investigation, we used the optical digital image correlation (DIC) technique [2] to obtain the deformation field on the specimen surface. This technique relies on the computer vision approach to extract the whole-field displacement data by comparing the features in a pair of digital images of a specimen surface before and after deformation.

An INSTRON 1125 screw-driven load frame loaded the specimens at a constant crosshead of 0.25 mm/minute. The applied tensile load and the crosshead displacement were monitored and recorded at a sampling rate of 10 Hz. A random speckle pattern was printed onto the specimen surface by first depositing a very thin white background and then spraying a black paint. A CCD camera, with the resolution of 1628 x 1236 pixels, was setup in front of the specimen. A series of images was captured during the test at the framing rate of 5 frames/second. All experiments were conducted in ambient temperature (21°C). For every sample configuration and loading angle, two tests were conducted, so there are total of 22 tests in this series of experiments.

Results and Discussion

Mode-I Loading

Figure 3 presents the mode-I (opening mode) strength measurements of all three types of specimens shown in Figure 2. The strength is defined, to be consistent for all configurations and loading angles, as the maximum applied load divided by the overall cross-section area of the specimen, but this definition does not account for the area of notches. Therefore the average interfacial strength for the notched specimens shown in Figures 2(a) and (c) will be twice the strength values presented in Figure 3.

Figure 3: Strength measurement of three different sample types subject to mode-I loading.

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The manner in which the specimen failed was also recorded by either camera images captured during the test or by inspection of the failed specimens after the experiment. For the Al-Al-notch specimens, subject to mode-I load, a very small amount of crack growth was observed followed by brittle failure along the Al/Al interface. Figure 4 presents the failure sequence of one of the Al-Al-notch specimens. Figure 4(a) shows the initial state of the specimen; Figure 4(b) is the moment prior to brittle failure where a very small amount of crack growth can be seen; and Figure 4(c) shows the sample after the instant of brittle failure. Because the framing rate of the camera is limited, only a fuzzy image of the sample is captured as the pieces flew apart.

Figure 4: Failure sequence of an Al-Al-notched specimen in mode-I loading: (a) initial image, (b) image of the specimen prior to the moment of failure, and (c) sample image after brittle failure.

The strength of the Al-DU-corner specimens, when subject to mode-I load, exhibits wide scattering, as shown in Figure 3. Note that the strength of one of the Al-DU-corner specimens is very close to that of the Al-Al-notch specimens, which indicates that the strength of this specimen was controlled by the Al/Al bonding alone and that the contribution of the bond between Al and the Zr/DU-10wt%Mo layer is negligible in this case. The other Al-DU-corner specimen exhibits much higher strength compared to any other mode-I specimens shown in Figure 3. For the weaker Al-DU-corner sample, interfacial cracks form along both the upper and the lower interface between the Al and the Zr/DU-10wt%Mo layer, and a crack also forms along the Al/Al seam in front of the corner of the Zr/DU-10wt%Mo layer. The stronger Al-DU-corner sample fails in a much more brittle and dynamic manner.

Finally, the strength of the Al-DU-notch specimens subject to mode-I loading is the lowest of all other specimens tested and the scatter is also large. The failure proceeds in brittle and dynamic fashion, and the current imaging system is limited in its framing rate, so that the details associated with the growing crack cannot be captured.

Mixed-Mode Loading

The strength of samples subject to mixed-mode loading, where the loading axis is at 45° to the interface of the CT specimen, are shown in Figure 5. Every type of specimen subject to mixed-mode loading is stronger than that when it is subject to mode-I loading, except for the one Al-DU-corner sample shown in Figure 3.

Al-Al-notch specimens clearly show an amount of crack growth before they fail. The Al-DU-corner specimens fail in a brittle, dynamic fashion without visible crack growth. When shearing

244

deformation was applied to the Al-DU-notch specimen, it was applied in both "positive" and "negative" shearing directions. By positive shearing, we mean that in this bi-material system, where the Al is sitting on top of the Zr/DU-10wt%Mo layer and an interfacial crack is situated along the interface, the shearing direction above the Al/Zr/DU-10wt%Mo interface in front of the notch tip is pointing away from the notch tip. If the shearing direction is pointing toward the notch tip, we call the situation negative shearing. As shown in Figure 5, the strength of the Al-DU-notch samples in both positive and negative shearing is essentially the same and they all exhibit brittle failure.

Figure 5: Strength measurement of three different sample types subject to mixed-mode loading.

Mode-II Loading

In mode-II loading, the external load is applied in the direction parallel to the Al/Al seam or the Al/Zr/DU-10wt%Mo interface, so that the notch-tip in the Al-Al-notch specimens and the Al-DU-notch specimens and the corner-tip in the Al-DU-comer specimens are initially subject to shear stress. The measured strength of all the samples is shown in Figure 6. The overall trend is that for all types of specimens subject to mode-II load, the strength is higher than that of both the mode-I load and the mixed-mode load as shown in Figures 3 and 5.

Figure 6: Strength measurement of three different sample types subject to mode-II loading.

When specimens were subject to mode II shear-dominated loading, both the maximum strength at which the sample fails and the appearance of failure of these samples are different than when when some portion of the loading is tensile. For the Al-Al-notch specimens in mode II loading

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(shearing mode), an amount of crack growth was observed during loading, but then the specimens did not fail. Instead, the bolt-holes gripping die specimen tore, so that the strength measurements shown in Figure 8 for the Al-Al-notch specimens only represent the value at which the pre-notch starts to propagate and the true strength value is some amount higher than this value. Similarly for the Al-DU-corner samples in mode II loading, neither the Al/Al seam nor the Al/Zr/DU-10wt%Mo interface failed, but instead the bolt holes fail. Therefore, we use the ">" symbol to indicate that the strength of the AI-DU-comer samples is no less than the values shown in Figure 6.

When the Al-DU-notch samples were subject to mode-II shearing in the negative sense, the situation is similar to that of the Al-Al-notch specimens, where a small amount of crack growth occurred along the Al/Zr/DU-10wt%Mo interface and the specimen never failed. The difference between Al-DU-notch samples and the Al-Al-notch samples when subject to negative shear is that we observe extensive deformation in front of the notch in the Al-DU-notch samples. On the other hand, when the Al-DU-notch samples were subject to the positive shearing load, extensive deformation in front of the notch was also observed, but this extensive deformation was followed by brittle failure. Thus the strength measurement shown in Figure 6 for Al-DU-notch sample subject to positive shear represents the true strength value.

The distinction of the failure behavior of Al-DU-notch sample subject to either negative or positive shear highlights the effect of the direction of shear at the tip of the notch when large deformation is involved. For infinitesimal deformation, the direction of the shear stress at the notch tip has no effect on the material failure. However, for finite or large deformation, even when the globally applied load is predominantly shear, the local deformation near the notch tip will involve some normal component. This normal component is compressive for negative shear, and it is tensile for positive shear.

Deformation Field and Related Results

In the preceding sections, the global strength measurement was presented and the observation of the appearance of the failure process was discussed. The strength measurement can be obtained from the global loading monitored by the load cell. The deformation or the displacement measurement is a challenging issue. Although we monitored and recorded the crosshead motion of the test machine, such a measurement does not represent the deformation or displacement experienced by the test specimen, since the compliance of the test machine and loading fixture also contributes to the motion of the crosshead. In this series of experiments, the displacement experienced by the CT specimen was determined using the optical DIC technique. Some of these results will be presented here.

DIC relies on the computer vision approach to extract the whole-field displacement data by comparing the features in a pair of digital images of a specimen surface before and after deformation. The features used in this series of experiments are the random speckle pattern on the sample surface. The random speckle pattern was made by first painting a thin layer of white background and then by spraying a black paint onto the surface. Figure 7 shows the displacement fields, obtained by using DIC, of the Al-Al-notch specimen subject to mixed-mode load very close to the moment of failure. The contour plot on the left is the displacement component in the

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horizontal direction and the contour plot on the right the displacement component in the vertical direction.

Figure 7: Displacement fields on the surface of Al-Al-notch specimen obtained from DIC.

One use of the DIC data shown in Figure 7 is to provide displacement information over a region on the surface of the specimen. One such region is illustrated in Figure 8(a) as the yellow dashed line. Among all the points from the DIC data, we can identify the displacements along the boundary lines on the top and the bottom of the region indicated as pink lines. The average relative motion of these two lines can be used as a measure of the displacement experienced by the CT specimen during the testing. Figure 8(b) presents the variation of the normalized applied load as function of the displacement determined according to the scheme just mentioned for two types of the CT specimens. One is the Al-Al-notch specimen subject to mixed-mode load and the other is the Al-DU-notch specimen subject also to the mixed-mode load with the positive shear component. The distinctive response of these two types of samples is apparent. The Al-DU-notch specimen, when subject to mixed-mode load, behaves in an elastic/brittle fashion, where the curves are linear up to the point of final failure.

Figure 8: (a) Extraction of the displacement experienced by the CT specimen from DIC data. (b) Variation of normalized applied load as function of displacement experienced by the Al-Al-notch

specimen and the Al-DU-notch specimen subject to mixed-mode-loading.

On the other hand, the Al-Al-notch specimen under mixed-mode loading resembles the behavior of elastic/plastic material, where following the apparent linearly elastic deformation, there is a very large portion of nonlinear deformation prior to the final failure. The contributing factor for such nonlinear deformation includes the plastic deformation near the tip of the notch and the small amount of crack growth along the Al/Al interface.

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The strain fields on the specimen surface can also be obtained based on the displacement fields obtained from DIC. Figure 9 shows the contour plots of the strain fields of the Al-Al-notch specimen subject to mixed-mode load close to the moment of failure. From the left to the right, these contour plots represent the normal strain in the horizontal direction, the normal strain in the vertical direction, and the shear strain. The quantitative measurement of the local deformation fields will provide useful information to FEA simulations to validate the numerical models.

Figure 9: Contour plots of the strain fields on the surface of Al-Al-notch specimen subject to mixed-mode loading.

Summary

A total of 22 compact tension (CT) tests were conducted. We investigated three different sample configurations, the Al-Al-notch, the Al-DU-corner, and the Al-DU-notch specimens. Three different loading angles were also evaluated: the mode-I (tensile opening) load, the mixed-mode load, and the mode-II (shearing) load. The DIC technique was applied to obtain full-field deformation contour maps of the specimen surface during the test. We observed that except one outlier in Al-DU-corner specimen, the strength of all the specimens increased when more shearing component was introduced in the loading. Subject to the mode-I (opening) load, the specimens (or the interfaces) exhibit the weakest strength and the widest scattering.

Acknowledgements

The authors would like to acknowledge the financial support of the US Department of Energy Global Threat Reduction Initiative Reactor Convert program. Los Alamos National Laboratory, an affirmative action equal opportunity employer, is operated by Los Alamos National Security, LLC, for the National Nuclear Security Administration (NNSA) of the U.S. Department of Energy under contract DE-AC52-06NA25396.

References

1. K.M. Liechti and Y.S. Chai, "Asymmetric Shielding in Interfacial Fracture Under Inplane Shear," J. Appl. Meek, 59(2) (1992), 295-304.

2. M.A. Sutton, et al., "Advances in two-dimensional and three-dimensional computer vision," Topics in Applied Physics, 77 (2000), 323-372.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

UNIAXIAL TENSION OF FRICTION-WELDED 304 STAINLESS STEEL AND 6061 ALUMINUM

C. Liu, M.L. Lovato, and W.R. Blumenthal

Materials Science and Technology Division Los Alamos National Laboratory, Los Alamos, NM 87545, USA

Keywords: 304 stainless steel, 6061 aluminum, DIC, Friction-welded, Tensile test

Abstract

A study of friction-welded 304 stainless steel and 6061 aluminum was conducted using uniaxial tension loading and optical two-dimensional digital image correlation (2D-DIC) to generate full field deformation maps on the surface of the specimen and to quantify the mechanical stress-strain response. A significant observation was that failure did not occur at the friction-welded interface, but instead occurred within a strain-localized neck region in the aluminum approximately 1.5 mm away from the weld interface. A detailed analysis of the 2D-DIC data set for one test specimen was used to illustrate how the mechanical response of the aluminum changed significantly from its "pre-weld" behavior due to the friction-weld process on a sub-millimeter length scale. A methodology for quantifying the mechanical property variations in the aluminum as a function of the distance from the welded interface was demonstrated using multiple, high spatial resolution, post-test "virtual extensometers".

Introduction

Uniaxial tension tests were conducted utilizing the two-dimensional digital image correlation technique (2D-DIC) [1] that demonstrated its powerful ability to quantify heterogeneous material response to sub-millimeter length scales. Tensile specimens were machined parallel to the normal of a plate of 304 stainless steel and 6061 aluminum bonded together by a friction-welding process that will not be discussed in this paper. Although sixteen tensile samples from three different welded plates were tested and analyzed, in this paper only one representative test will be used to demonstrate a methodology for quantifying the mechanical property variations in the sample using multiple, high spatial resolution, post-test "virtual extensometers".

Axi-symmetric "dog-bone" geometry specimens were machined with a gage diameter of 1.6 mm and a nominal gage length (straight section length) of 12.7 mm, so that the friction-welded interface was located in the center of the tensile specimen and perpendicular to the loading axis. A 1 kN load cell was used and the maximum tensile loads were approximately 500 N. Tests were conducted at 20°C and the specimens were loaded at a constant cross-head speed of 0.66 mm/minute using an Instron model 1125 screw-driven testing machine. The nominal plastic strain rate for a homogeneously deforming specimen would be about 8 x 10"4 per second; however, as will be revealed later, the deformation response of these composite samples varied from purely elastic (in the 304 stainless steel) to highly localized within a "neck" region in the 6061 aluminum. Hence, the plastic strain rate was not a constant, but varied from zero to many times the nominal strain rate within the necked region.

One CCD camera with a resolution of 1628 x 1236 pixels was setup in front of the specimen to record digital images with two banks of LED lights illuminating the specimen from both left and

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right sides. A random speckle pattern was applied to the specimen surface by first depositing a very thin white background and then mist-spraying fine spots with black paint. The image acquisition rate was 5 frames/second and each test ran for about 1 minute. Thus a total of about 300 images were generated and analyzed via D1C for each test. The physical resolution of the random speckle image is approximately 11 micrometers/pixel.

When processing the random speckle images using the DIC technique, the ROI (region-of-interest) was identified using image processing on the reference image and it was chosen to be slightly larger than the gage length to allow for capturing considerable tensile elongations. The spacing between DIC data points is set at 5 pixels and the subcell size is fixed at 21 pixels, which is equivalent to a strain gage rosette with gage size of 0.23 mm. Incremental correlation (compared to an initial reference image) is used so that there is no need for specifying the "initial guess" of each frame. Since the deformation in the 6061 aluminum involves large plastic strains, the logarithmic (true) strain was calculated and the "window size" (the number of data points for calculating the derivative) was chosen to be 15.

Uniaxial Tensile Testing and Analysis of Specimen 2135-3

The 2D-DIC measurements and subsequent analysis scheme generate a large amount of data (over 300 MB per test), so we present in this paper a detailed analysis for only one representative test (specimen number 3 of plate 2135) to demonstrate the capabilities of this technique and a methodology for quantifying the highly localized mechanical response of the friction-welded composite plate materials near the interface.

Figure 1 : Overall response of uniaxial tension sample 2135-3.

The overall engineering stress-strain response of tensile specimen 2135-3 is shown in Figure 1 as a reference for subsequent 2D-DIC deformation analysis. In this plot, the horizontal axis is the engineering tensile strain (e.g. the crosshead displacement, A, normalized by the initial length of the straight section of the specimen, L) and the vertical axis is the engineering tensile stress (i.e. the applied tensile load, P, normalized by the initial cross-section area of the specimen, A = n D214, where D is the initial gage diameter of the sample).

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During the test, the engineering tensile stress initially increases with nominal strain in an almost linear manner and then exhibits apparent yielding near Point A; the stress then continues to increase until it reaches a maximum at Point B; the stress then gradually softens (through Point C); and then drops rapidly beyond Point D. Point E is the last moment that the tensile specimen is still held together, after which the sample completely fails and breaks into two pieces.

Local displacement fields and strain fields were analyzed using the 2D-DIC technique and are presented as color contour plots overlaid on video images of the deformed specimen in Figure 2 at each of the Points A-E. Figures 2a and 2b show contour plots of the axial displacement field, u, and the axial tensile strain field, ex, respectively. The evolution of a complex, heterogeneous deformation field during the tensile loading process is apparent on the aluminum side (right half) of the interface. Prior to the maximum load at Point B, local deformation in both the stainless steel and the aluminum sides are very uniform, although the relative deformation within the stainless steel is proportionally smaller than in the aluminum due to the substantial stiffness mismatch of the two materials. Then at the moment of maximum load (B), deformation localization within the aluminum becomes very noticeable and a "neck" is formed. As the test continues through Point C to Point E, strain localization in the neck region intensifies, but at a distance from the friction-welded interface at the center of the specimen. Finally, the specimen fails and breaks near the middle of the necking region just after Point E.

Figure 2: Contour plots of (a) axial displacement field, «, and (b) axial tensile strain field, e ,̂ at the selected points, A-E, indicated in Figure 1.

By manipulating the DIC data, any number of "virtual extensometers" can be established at different axial locations on the specimen, analogous to physical extensometers, to extract local axial tensile displacements and strains as a function of the local stress. Furthermore, virtual

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extensometers can also be specifically tailored to different length scales at key locations, such as the neck region, where local strains and stresses are the greatest and represent the full extent of a material's tensile response to failure.

Figure 3: Virtual extensometers used for measuring deformation in the aluminum portion of the uniaxial tension specimen 2135-3.

The contour plots of the displacement field, u, and the strain field, er, shown in Figure 2, indicate that all of the plastic deformation occurs in the aluminum portion of the tensile specimen (as might be expected). Thus, further analysis of the deformation of the aluminum side was performed by applying four different virtual extensometers as shown in Figure 3, all of which capture behavior in the localized neck region. Three of the extensometers have finite gage lengths of 6 mm, 4 mm, and 2 mm, respectively, with one end of each extensometer fixed at the friction-weld interface that separates the stainless steel and the aluminum (indicated by the dotted green line in Fig. 3). The axial strains from these three gages are computed from the relative extension of the two ends of each gage using the following definitions. The axial true (or logarithmic) strain is given by e = In(Z./L0), where L is the current gage length and L0 is the initial gage length. The axial true tensile (or Cauchy) stress, c, is defined as the applied tensile load, P, normalized by the current cross-section area. Assuming the material is incompressible, the true strain can be used, such that o = (P I{KD2 /4))exp(-e), where D is the initial gage diameter. Gage Number 4 is special because it has no effective gage length and is located where final necking and failure occurred. It measures the mean value of the local strains along a single cross-section of data obtained from the DIC measurements perpendicular to the loading axis.

The true stress versus true strain for all four virtual gages in the aluminum is plotted in Figure 4. Clearly there are large differences in the apparent ductility measured by the different virtual extensometers. The shorter extensometers are more dominated by the strain localization in the neck region and therefore display higher ductility because the deformation is not homogeneous along the axis, but is highly localized. Because Gage Number 4 is located right at the peak of strain localization in the neck, it captures the maximum axial strain (ductility) associated with failure of the sample. More importantly, after yielding at approximately 180 MPa, the hardening behavior measured by the different virtual extensometers is not the same curves when it should be identical for a homogeneous material. Hence the aluminum near the friction-welded interface of the samples has become significantly inhomogeneous due to the highly dynamic thermo-mechanical nature of the friction-welding process. This hypothesis was confirmed by testing a homogeneous 6061 aluminum specimen that was fully annealed (to an over-aged condition).

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Figure 4: Axial tensile true stress, a, as a function of axial true strain, e, of the aluminum portion of sample 2135-3 measured with different virtual extensor-meters shown in Fig. 3. The initial gage lengths are in parentheses.

As shown in Figure 5, the strain-hardening response of annealed 6061 aluminum is independent of location along the axis of the specimen, verify that the specimen is homogeneous.

Figure 5: Tensile stress, c, versus tensile strain, e, of a homogeneous aluminum specimen using virtual extensometers at different locations shown in the inset.

The investigation will now focus on attempting to quantify the variations in the mechanical properties of the friction-welded aluminum at sub-millimeter spatial resolution as a function of the distance from the welded interface.

The same DIC strain measurement technique used for Gage Number 4 in Figs. 3 and 4 (at a single cross section) was applied at five additional sites in the aluminum section of specimen 2135-3 as indicated in Figure 6. The initial distances of each site to the stainless steel/aluminum weld interface were: 0.569 mm, 1.082 mm, 1.594 mm, 2.106 mm, 2.619 mm, and 3.074mm,

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respectively. Site 3 is the same as Gage Number 4 in Fig. 3 and again is the location in the neck where final failure occurred. The true tensile stress versus true tensile strain at each of the six sites is plotted in Figure 7.

Figure 6: Six cross section sites in the aluminum portion of specimen 2135-3 where deformation was monitored relative to the failure location (Site 3).

Figure 7: Tensile true stress, a, versus true strain, e, is plotted for six cross-sections in the aluminum portion of sample 2135-3. Initial distances to the friction-weld interface are in parenthesis. The inset magnifies the low strain region.

Except for Site 3 (the failure location), the curves at the other sites are only plotted up to their maximum stress because the stress actually unloads in these regions after this point as necking develops. Hence, only Site 3 exhibits a "full" stress-strain curve to failure, while the other sites exhibit truncated responses. The inset plot within Figure 7 is a magnification of the low strain region that better highlights the differences in the strain-hardening response at each site. Plainly, the aluminum at Site 3 is the "softest" among all the locations and this is the reason necking and eventual failure occurs here. Note that Sites 2 and 4 and Sites 1 and 5 are equidistant pairs from Site 3. However, the sites closer to the weld interface on the left of Site 3 are softer than the corresponding pair sites on the right side of Site 3. Site 6 is not paired, but is the farthest of the selected sites from the weld interface and is perhaps most representative of the "pre-welded" aluminum. The key observation from Fig. 7 is that the "pre-welded" aluminum has undergone

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significant changes within very short distances of the weld interface and is now very heterogeneous, but not in a monotonie manner, resulting from the friction-weld process.

A limitation regarding 2D-DIC strain measurements, which uses only a single camera, needs to be pointed out here. During uniaxial tensile testing of specimens with cylindrical cross sections, the initial motion of points on the curved sample surface is along the direction parallel to the loading axis. Since the loading axis is aligned parallel to the image plane of the CCD camera, this motion is a true representation of the axial tensile strain. However, when a neck develops, then the sample surfaces on the "shoulders" of the neck are no longer parallel to the image plane and strains measured at those locations after necking develops should be regarded as approximate. This caveat only applies to Sites 2 and 4 in Fig. 7 and does not alter the observation that the aluminum undergoes significant changes due to the friction-welding process.

the various sites, a simple model is used that

for e < EY

( < ) for £ > Ey

where £Y and aY are the yield strain and the yield stress, respectively. Parameter N is the strain hardening exponent and is within a range 0 < N < 1. N = 0 corresponds to elastic-perfectly-plastic response. We will apply Eq. (1) to the curves shown in Fig. 7 to determine the material characterizing parameters Ey, aY, and N at the different site locations.

The yield stress and strain point for each site was determined using the standard 0.2% off-set method where the linear-elastic portion of the stress-strain curve is fit to a straight line and is then offset-shifted by 0.2% strain to find the intersection point on the stress-strain curve that defines the yield point, (er,ar).

Note that the power-law representation of the strain-hardening exponent N, in Eq. (1) is unable to properly represent the post-yielding behavior where the true stress reaches a maximum and the slope of the stress-strain curve drops to zero. Therefore the strain-hardening exponent, N, was only fit to the data close to the yield point for each curve using explicit Eq. (2):

-felts)- - -where the slope of the stress-strain curve, da I de, is determined by fitting a straight line to the three or four 2D-DIC data points on either side of the yield point, (er ,or ).

Material parameters, er, ay, and N, determined at the six different sites in the aluminum are plotted in Figure 8 as function of the distance from the weld interface. The yield strain er

measured at different sites ranges from 0.29% to 0.44%, where the lowest value (0.29%) was at Site 3. The yield stress, a,, also has its lowest value at Site 3 where necking and final failure occurred. The yield stresses measured from Sites 2 and 4 are slightly higher than that at Site 3, but at Sites 5 and 6, the yield stress is markedly greater. The strain hardening exponent, N, has a

To quantify the variations in material properties at relates the true stress, a, and the true strain, e,

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V-shape as a function of distance from the interface with a minimum at Site 3 and with exponents relatively higher on the right-hand side of Site 3, farther from the weld interface.

Figure 8: Material parameters at different cross section sites in the aluminum portion of specimen 2135-3 as a function of distance from the weld interface.

Conclusions

A study of friction-welded 304 stainless steel and 6061 aluminum was conducted using uniaxial tension loading and two-dimensional digital image correlation (2D-DIC) for measuring full-field displacements and strains on the specimen surface. Several interesting observation were made about the deformation and failure response. First, failure did not occur at the friction-welded interface, indicating a relatively strong bond, but instead occurred in the aluminum section approximately 1.5 mm away from the weld interface and within a highly strain-localized neck region. Second, the 6061 aluminum underwent significant changes from its "pre-welded" condition due to the friction-welding process. Instead of remaining homogeneous, the aluminum becomes very heterogeneous in a relatively complex manner and within a very short distance of the weld interface. The 2D-DIC technique and a methodology for quantifying the mechanical property variations in the aluminum on a sub-millimeter length scale using multiple, high spatial resolution, post-test "virtual extensometers" was successfully demonstrated.

References

1. M.A. Sutton, et al., "Advances in two-dimensional and three-dimensional computer vision," Topics in Applied Physics, 11 (2000), 323-372.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ULTRA FAST GRAIN BOUNDARY SEGREGATION IN HOT DEFORMED NICKEL

Marion Allart, Frederic Christien, René Le Gall

University of Nantes, Polytech Nantes - LGMPA Rue C Pauc, BP 50609 44306 Nantes Cedex 3, France

Keywords: Nickel, sulfur, WDS, Wavelength Dispersive X-ray Spectroscopy, diffusion, intergranular segregation, grain boundary, deformation, subgrain

Abstract

Hot deformation of nickel alloys (500-800°C) can lead to a dramatic loss of mechanical properties due to sulfur grain boundary segregation. In this work, both equilibrium and non-equilibrium sulfur grain boundary segregation in nickel were studied at 550°C. They were obtained respectively by simple annealing and by hot-compression. Segregation was quantified by Wavelength Dispersive X-ray Spectroscopy and was found to be thousands of times as fast during hot-compression as during annealing at the same temperature: it takes about 20 minutes at 550°C to saturate the grain boundaries with sulfur during compression at 3.10'5 s"', whereas three months are needed by simple annealing at the same temperature. Cross-section observations show that subgrains form during hot compression. A model based on sulfur diffusion along the subgrain boundaries (dislocation walls) enables to account for the observed ultra-fast segregation during hot-compression.

Introduction

Alloys manufacturers observed cracks on nickel ingots during hot deformation. This damaging has been attributed to sulfur intergranular segregation. Interfacial segregation of impurities in metals is known to possibly cause tremendous changes in the materials physical properties. Among metal-impurities couples, one of the most studied is nickel-sulfur. Its thermodynamic and kinetics parameters have been widely studied. However, during hot-deformation, non equilibrium segregation occurs at a very high speed that literature cannot explain. It's thus obvious that there is a mechanism of segregation to identify. From this perspective, we decided to compare the sulfur segregation between annealed and hot-deformed nickel at a temperature of 550°C. The first part of this paper will deal with the technical part of the study. We will describe the metal we used and the way segregation was obtained, by annealing and by hot compression. The quantification of segregation thanks to Wavelength Dispersive X-ray Spectroscopy (WDS) will also be exposed in this chapter. In the second part, we will state the results obtained on both annealed and deformed samples and compare them. The assumed mechanism involved in deformation-accelerated segregation is proposed at the end of the paper.

Material And Techniques

Material

The material we used has been specially fabricated by ArcelorMittal-Aperam for this study with a very low concentration of sulfur in order to avoid the precipitation of nickel sulfur. An analysis by Glow-Discharge Mass Spectrometry (GDMS) showed a concentration of sulfur of 5.4 wt ppm

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(9.9 at ppm). As we need a uniform and desegregated initial state, the whole metal has been annealed 30 min at 900°C in argon. The nickel is cut into 6x3x20 mm pieces. Each sample is annealed once more 30 min at 950°C before every experiment. This last heat treatment is made under vacuum to avoid oxidation.

Annealing

As mentioned above, the first results deal with samples annealed at 550°C. Each one is sealed in a tube of quartz under vacuum to avoid nickel oxidation. The sealed tubes are placed in a furnace at the temperature of 550°C for 1 to 100 days.

Hot Compression

We want to replicate the phenomenon of accelerated segregation occurring during hot deformation. The Figure 1 represents the system used to heat and compress the samples. The sample is placed in a tensile machine. An arc-image furnace, specially designed for this machine, surrounds the sample. The temperature is controlled via a thermocouple located in the upper part, as close to the sample as possible.

Figure 1. Schematization of the system used to heat and compress samples

Hot compression process is automated to enhance precision and reproducibility. The compression involves four steps:

1 / Heating up to 550°C while a slight pressure (20 N) is maintain on the sample to ensure thermal contact between the supports and the sample (15-20 min); 2 / Wait until the thermal equilibrium is reached (50 min); 3 / Compression at a given speed until the expected deformation (1-20 min); 4 / Quenching of the sample in air to ensure a rapid cooling to stop segregation.

During the program the software acquires the strength, the displacement and the temperature. It should be noted that a high concentration of dislocations, induced by the deformation, can lead to dynamic recrystallization. This phenomenon is not the one we want to study and it must be avoided. One of the consequences of dynamic recrystallization is a much smaller grain size. This is obvious when observing the fracture surface in a SEM and we can easily exclude from our results any sample that was subjected to dynamic recrystallization.

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Wavelength Dispersive X-rav Spectroscopy

The quantification of segregated sulfur in the grain boundaries is obtained by Wavelength Dispersive X-ray Spectroscopy (WDS). This technique has been used for decades, but its application to the measure of mono-atomic layers of segregation is a recent improvement made in our laboratory. We will give here a quick presentation of this technique, but the reader can find a detailed version in the published papers [1,2].

Application to measuring grain boundary segregation As we use a surface tool to study an interface, the first step is to obtain a brittle and intergranular rupture of our material, which is made by cryogenic traction in liquid nitrogen. This rupture of the sample is made ex situ as the WDS is not sensitive to contamination from the atmosphere. This is an important advantage of this technique with respect to Auger Electron Spectroscopy, where rupture must be obtain in situ in ultra-vacuum. It must be specified that not every grain boundary can be observed this way. The rupture will be intergranular if the concentration of sulfur in the grain boundary is at least 15 at%, equivalent to approximately 15 ng.cm"2 [3]: only grain boundaries with a concentration of sulfur superior to 15 ng.cm"2 can be studied this way. The fractured surface is uneven, as can be seen on Figure 2, which affects the WDS measurements; the tilt angle 6 has to be taken into account in the quantification of segregation. Our team developed an original method to obtain the tilt angle between the electron beam and the facet being analyzed by using the value of the sample current. Depending on the tilt angle, the quantity of secondary and backscattered electrons will vary and so will the absorbed current, or sample current, that can be measured. The equation of quantification of the sulfur grain boundary concentration u, applied to intergranular segregation, is:

/ I C" -C (1) M = 2K—— x cosö = 2K—— x i ^ - ^

I I C -C 1 Stamford 1 Standard ^ a b s ^'b

With K Parameter depending on the element measured, the bulk material, the standard material used and the beam voltage; in our case K. = 256.1 gem"2

/ Intensity of the peak measured on the sample [count.s"'.nA"'] htandarä Intensity of the peak measured on the standard material [counts"'.nA'1]

C°abl Absorbed current measured on a nickel sample normal to the beam [nA] CB Absorbed current measured on the facet being analyzed [nA]

The result of the WDS quantification is expressed as a mass of sulfur per unit surface (ng.cm2). This can be converted to a number of S atoms or a fraction of a monolayer. For example, assuming that a monolayer is a (111) Ni plan, 50.0ng.cm"2 of sulfur corresponds to 9.39.10''atoms per cm and to 50.4 % of a monolayer.

Experimental conditions The scanning electron microscope used for these measurements is a Merlin SEM from Carl Zeiss, equipped with a field emission gun. The WDS measurements were carried out on a INCAWave 500 spectrometer from Oxford Instrument. The monochromator crystal used is a PET (Pentaerythriol). In this work, all the analyses were carried out at 20 kV. Quantification of sulfur is based on the measurement of the Ka line of the sulfur. It's theoretical position is 5.3730 Â (2.308 keV), but a scan made before every experiment checks its real position. As scanning the whole peak would be needlessly time-consuming, three measures only are made on each analyzed zone, one on the peak position, one on the left (5.1999 Â), and one on the right (5.5500 Â). The intensity of the peak is deduced assuming a linear background.

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The standard used for the quantification is FeS2. The beam voltage used on the standard is 20 nA, to avoid saturation. Photons were counted 60 s on the peak and 30 s on each background. On the samples, the concentrations of sulfur are very low, so optimization was made by maximizing the beam with a current of 450 nA and counting 180 s on the peak and 90 s on each background. In order to improve statistics, 12 to 20 measurements by scanning 15x20 urn zones are made on each sample, one measure corresponding to one facet. On Figure 2, the areas analyzed by WDS are represented by black squares.

Figure 2. SEM image of a fracture surface of a sample annealed 100 days at 550°C. Each black square represents a zone analyzed by WDS.

Experimental Results

Equilibrium segregation at 550°C

Figure 3. WDS quantification of the concentration of sulfur segregated to the grain boundaries on nickel samples annealed at 550°C ( ■ ) and deformed at 550°C ( ► ) (a) log time scale (b) square toot time scale

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Experimental results obtained on samples annealed at 550°Care represented by the square dots on Figure 3. The sulfur surface concentration measured on the grain boundaries is plotted with respect to the annealing time. The error bars is the standard deviation calculated over the measurements made on the same sample. The time-dependence of the sulfur grain boundaries concentration is in square-root (Figure 3(b)), which is consistent with the well-known McLean approach of interface segregation kinetics. The value of the concentration of sulfur in the grain boundaries in the initial state appears to be approximately 15 ng.cm'2. This can be attributed to two concomitant phenomena. The first one is that we can only quantify brittle grain boundaries, that is to say with more than 15 ng.cm"2 sulfur. This restriction makes it impossible de facto to quantify very low levels of segregation, such as samples submitted to short annealing times (<20 h). It also can lead to a slight overestimation of segregation as we can only make measurements on the grain boundaries with the highest level of segregation. The second one is that despite the thermal treatment at 950°C, the nickel grain boundaries we presumably not totally desegregated prior to the annealing at 550°C. After long-time annealing at 550°C, the intergranular sulfur concentration should be stable. As we do not see a plateau on our results (Figure 3), it means that 100 days is not enough to reach final equilibrium.

Segregation during deformation at 550°C and 3.10 s s"1

Samples of nickel were compressed at 550°C at a deformation rate of 3.10"5 s'1. The kinetics of grain boundaries sulfur segregation during hot-deformation is plotted with black triangles on the Figure 3. The bottom x-axis indicates the time of compression, the top one indicates the deformation e calculated from the initial and final height (h0 and hj) of the sample:

The striking result is the small amplitude of deformation needed to have a significant effect on the intergranular segregation: a deformation during 20 min, until a deformation of 4%, leads to a grain boundaries concentration of about 50 ng.cm" , which is the same effect as a 90 days of annealing without deformation!

Microstructure

In order to compare the microstructure of the nickel before and after hot compression, samples were prepared by ionic polishing using a Cross Section Polisher from Jeol. The observations were made by backscattered electron imaging. The images are presented on the Figure 4.

261

Figure 4. Backscattered images of samples of nickel before (left column) and after (right column) a hot compression at 550°C, SJO'̂ s ' to a deformation of E = 1.81 %.

On the backscattered electrons images of the nickel before deformation, each grain is uniform. There is no crystallographic misorientation, at any scale, as can be seen on images (a) and (b). On the contrary, on the hot-compressed samples (image (c)), a crystalline contrast is visible on each grain. At high magnification, like on image (d), the structure appears to be a juxtaposition of micrometric zones of uniform orientation (subgrains) separated with very thin zones of high misorientation, in white on the picture. This microstructure is the result of the dislocations rearrangement. The dislocations created by the deformation tend to move until reaching a grain boundary or getting snarled up with some other dislocation [4]. They stack up and form dislocation walls: the white stripes on image (d).

Discussion

Equilibrium segregation at 550 C

As mentioned earlier, the kinetics of equilibrium segregation can be described by the McLean's simplified Law [5]:

, , c .VÖ7 (2)

f>-M + 4 — r -

Va-

in which pi is the surface concentration of sulfur in the grain boundary after / seconds in a

material containing a bulk concentration of sulfur Cv, fto is the initial concentration of sulfur in

the grain boundary and Db is the bulk diffusion coefficient of sulfur in nickel.

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Our experimental data were fitted with this law, as represented by the straight line on Figure l.The best fit was for a diffusion coefficient Dt, of 1.89.10"14 cm2.s"'. The excellent agreement with the value from the literature [6], which is 1.86.10"14 cm2.s"' at 550°C, definitively validates our protocol of quantification of sulfur segregation in the nickel grain boundaries.

Segregation during deformation at 550 C and 3.10" s"1

The same model used for equilibrium segregation can be applied to segregation during deformation, to obtain the sulfur "apparent" diffusion coefficient in nickel during deformation Dde/. The fitted line is the dotted line on Figure l.The result is a coefficient D^j of 6.03.10'" cm2.s"'. In other words, the segregation occurs more than 3000 times faster than in the case of simple annealing. We showed that dislocation walls form in the nickel dunng hot compression. This can explain the acceleration of sulfur intergranular segregation during deformation, as the structure of those walls is very close to the structure of grain boundaries. Dislocations mobility increases with temperature. In our case, we can suppose that the dislocation walls form quasi immediately, in the first times of deformation, as represented on the left of the Figure 5. It's also reasonable to postulate that sulfur atoms are in the walls as soon as they form. Indeed, the distances are quite short as the size of the dislocation cells is micrometric, and the migration of dislocations toward the walls, associated with a high mobility of the sulftir due to the temperature, strongly favors the sulftir segregation to the walls, by dislocation-dragging of sulfur atoms.

Figure 5. Schematization of sulfur segregation to grain boundaries by diffusion along dislocation walls

The dislocation walls have a structure close to grain boundaries and the sulfur diffuses really faster in the grain boundaries than in the bulk. We can assume that the sulfur will diffuse very fast along the walls. In other words, dislocation walls act like diffusion short-cuts. The experimental sulfur diffusion coefficient in the hot-deformed nickel found above is 6.03.10 " cm2.s'. It's very high when compared to the coefficient for bulk diffusion (1.89.10"'4 cm2.s"') but unexpectedly lower than the grain boundary diffusion coefficient that is 4.54.10-9cm2.s"'[7].

Conclusion

Our study aimed to identify the mechanism of ultra fast intergranular segregation of sulfur in nickel during deformation at 550°C. We measured the sulfur grain boundaries segregation kietics in annealed samples and in hot-compressed samples This quantification was made by WDS applied to intergranular segregation, an original technique developed in our laboratory. We measured a bulk sulfur diffusion coefficient in nickel in the annealing samples of 1.89.10'14 cm2.s ', in excellent agreement with the literature value of 1.86.10"14 cm2.s"'. In the case of hot-deformed nickel, a fit of the kinetics gave a diffusion coefficient of 6.03.10 ' cm .s1.

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In other words, the segregation occurs more than 3000 times faster than in the case of simple annealing. Those results tend to validate the hypothesis stating that the ultra-fast segregation is due to dislocation walls formed during hot-deformation and acting like diffusion short-cuts.

References

1. F. Christien, R. Le Gall, "Measuring grain boundary segregation using Wavelength Dispersive X-ray Spectroscopy", Surface Science, 602 (2008), 2463-2472.

2. P. Nowakovski, F. Christien, M. Allart, Y. Borjon-Piron, R. Le Gall, "Measuring grain boundary segregation using Wavelength Dispersive X-ray Spectroscopy: further developments", Surface Science, 605 (2011), 848-858.

3. JH. Heuer, P.R. Okamoto, N.Q. Lam, J.F. Stubbins, "Relationship between segregation-induced intergranular fracture and melting in the nickel-sulfur system", Applied Physical Letter, 76 (2000), 3403-3405.

4. F. Christien, R. Le Gall, G. Saindrenan, "Application of percolation theory to surface segregation during recovery", Ada Materialia, 51 (2003), 521-534.

5. D. McLean, Grain boundaries in metals (Clarendon Press, Oxford, 1957).

6. A. Vladimirov, V.N. Kaigorodov, S.M. Klotsman, I.S. Trakhtenberg, Fizika metallov i metallovedenie, 39 (1975), 319.

7. R. Le Gall, E. Quérard, G. Saindrenan, H. Mourton, D. Roptin, "Diffusion of sulfur from the grain boundaries to the surface of polycrystalline nickel", Scripta Materialia, 35 (1996), 1175-1181.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

QUANTITATIVE NANOSIMS ANALYSIS OF GRAIN BOUNDARY SEGREGATION IN BULK SAMPLES

F. Christien1,2, C. Downing1, K.L. Moore1, C.R.M. Grovenor1

1 Department of Materials, University of Oxford, Parks Road, Oxford OX I 3PH, United Kingdom

2 LGMPA, Université de Nantes, Polytech'Nantes, Rue Christian Pauc, 44306 Nantes Cedex 3, France

Keywords: SIMS, microbeam techniques. Auger Electron Spectroscopy (AES), traces, surface analysis.

Abstract

Solute grain boundary (GB) segregation is an important metallurgical phenomenon that has been extensively studied over the last 40 years, especially by Auger spectroscopy of fractured surfaces. More recently, it has been demonstrated that High Resolution SIMS (NanoSIMS) analysis can detect solute GB segregation on a simple polished cross-section. The aim of the work presented here is to demonstrate the use of SIMS to achieve quantitative analysis of GB segregation.

Introduction

Segregation of impurity elements to the grain boundaries in metals can lead to dramatic degradation of their mechanical properties. For example sulfur is known to strongly embrittle nickel alloys [1,2]. Grain boundary segregation has been extensively studied over the last 40 years, mainly using Auger Electron Spectroscopy (AES) [3]. The conventional route for AES analysis of grain boundaries is to fracture the sample in ultra-high vacuum and then to analyze the fracture surface. As most solute segregation phenomena of practical and industrial interest are associated with grain boundary cmbrittlement, the sample breaks in an intergranular fashion, allowing the analysis of the grain boundary chemistry by studying the exposed surfaces. It has also been shown that secondary ion mass spectroscopy can be used to detect grain boundary segregation on a simple polished cross-section [4-8], with the advantage that no sample fracture is required so that all grain boundaries can be studied not just those that are embrittled. Nevertheless, only qualitative data on grain boundary segregation by SIMS has been published so far. The aim of this work is to demonstrate that the NanoSIMS can quantitatively determine grain boundary equilibrium segregation levels.

Methods and Materials

The material used for this work is a 3 mm thick pure nickel sheet provided by ArcelorMittal Aperam (Imphy, France) with nominal bulk sulfur content 5.4wtppm (measured by Glow Discharge Mass Spectroscopy). Samples ( 3x5x10 mm) were cut from the sheet and annealed under high vacuum at six temperatures (550, 650, 750, 900, 950 and 1000°C) to reach different levels of segregation. Two different samples were annealed at 550 C and one sample at the other temperatures. To ensure that equilibrium segregation was achieved,

265

annealing times were calculated using the McLean kinetics equation [9] and the sulfur bulk diffusion coefficient in nickel from [10]. After annealing, the samples were mechanically polished down to 1 urn, cleaned with ethanol and introduced into the NanoSIMS. Standard materials As the aim of this study is to achieve quantification, standard materials of known composition are needed. Four pure nickel standards with different sulfur concentrations were used (0.5, 38, 73 and 134 wt ppm respectively). These materials were provided by ArcelorMittal-Aperam (Imphy, France) as 1 kg ingots (except the 0.5 ppm sample which is from Wiggins Alloy). These standard materials were annealed briefly at 1250°C and water-quenched to achieve a supersaturated solid solution. The sulfur concentration was then checked in each standard sample using Electron Probe MicroAnalysis (EPMA), and very good agreement was found between the EPMA measurements and the sulfur concentration specified by the provider.

Results and Discussion

SIMS analysis was performed using the CAMECA NanoSIMS 50 ion microprobe which allows simultaneous detection of five ionic species from the same sputtered volume with high mass and high spatial resolution while still maintaining up to 50°o transmission of all the secondary ions which is necessary to achieve high sensitivity. A focused 16 keV Cs+ primary ion beam is scanned over the surface of the sample and the sputtered negative secondary ions are collected and analyzed using a double focusing mass spectrometer. The primary Cs+ beam current was about 1.5 to 2 pA for all experiments. Calibration of the 32S"yield against bulk samples of known S content has been described elsewhere [11]. About 15 large areas (150x150 urn) on each sample were automatically cleaned overnight for one hour each using the maximum primary ion beam current, implanting a dose of about 1016 Cs+ ions cm" . Figure 1 shows a 20x20 um area around a triple point in a sample annealed at 550°C. Sulfur grain boundary segregation is clearly demonstrated on the 32S" map. It is also quite obvious that the sulfur segregation level is not the same in the three grain boundaries.

Figure 1 : (a) Secondary electron image of a 20x20 urn area around a triple point in a nickel sample annealed at 550°C. (b) Corresponding 32c-

S map.

Quantification of grain boundary segregation was achieved from 205x205 pixel sulfur maps acquired on 4x4 urn areas containing a region of straight grain boundary with a counting time of 0.015 s pixel. About 20 grain boundaries were analyzed for each sample. Figure 2 shows a typical set of 32S" maps and resulting quantitative linescans obtained from the seven samples.

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The calculated grain boundary apparent concentration, Q, is indicated on each profile. The values of Q can be determined from the peak area measured from the concentration profile. We have estimated the limit of detection of S at a single Ni grain boundary to be -0.24 ng cm"2 which corresponds to 0.004 of a monolayer [11], demonstrating the very good sensitivity that can be achieved in this combination of segregant and matrix. The real sulfur grain boundary concentration fi in g cm"2, taking into account the inclination of the grain boundary to the vertical plane, 0, can be calculated from [11]:

// /2cos<9 (1)

Because there is no simple method to measure 6 for each analyzed grain boundary we have estimated in a polycrystalline sample with equiaxed grains that the average 0 angle over a large number of grain boundaries is about 33° resulting in an average cosöterm of 0.79 [11] so that:

M 0.79/3 (2)

Figure 3 shows the distributions of the measured sulfur apparent concentration Q over the analyzed grain boundaries on three samples with different annealing temperatures (550, 750 and 950°C). The rather wide distribution of the apparent sulfur concentration in a single sample can be accounted for by the variation of the 6 angle as well as the variation of the real sulfur concentration from one grain boundary to another. In order to extract the thermodynamics of the segregation phenomenon we have averaged the sulfur grain boundary concentration over all the grain boundaries analyzed in each material. For the sample annealed at 950 C, for instance, this results in a value of fi 2.4 ng cm"2, which corresponds to an average concentration of 4.5x1013 sulfur atoms per cm2. This can be converted into a fractional monolayer coverage assuming that the atom density of a monolayer is the same as that of a (110) nickel plane (1.14x10 atoms cm"2) to give an average grain boundary segregation at this temperature of 0.04 monolayers of S. Figure 4 shows the average sulfur grain boundary concentration determined from the NanoSlMS analyses on each sample. As expected, the equilibrium sulfur grain boundary concentration decreases with annealing temperature. We have analyzed these data using the well-known Langmuir McLean formalism and obtained a free energy of equilibrium segregation, AG, of -97.8 kJ mol"1 and a saturation level of fiMax ~ 28.0 ng cm 2, which corresponds to a saturation concentration of sulfur atoms to an average grain boundary Nuax

5.27xl014 atoms cm"2. The AG value determined in this work is in excellent agreement with the value estimated by Larère from AES analysis (-98 kJ mol"1 at 750 C [1]), and the saturation concentration can be converted to 0.46 of a monolayer in very good agreement with values determined by Larère (XMax - 0.44 [1]) and by Lejcek (XMax 0.40 [12]).

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Figure 2: Typical grain boundary S' maps and corresponding concentration profiles acquired on the samples annealed at 550°C, 650°C, 750°C, 900°C, 950°C and 1000°C. The grain boundary apparent concentration i2(peak area) is indicated on each profile.

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Figure 3: Distribution of the sulfur grain boundary apparent concentration, Q, measured by NanoSlMS on the samples annealed at 550 C, 750°C and 950 C. For the sample annealed at 950°C, the small gray bar at 12 ~0 corresponds to boundaries where no excess sulfur was detected.

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Figure 4: NanoSIMS measurements of sulfur grain boundary segregation in nickel. The line is the Langmuir-McLean expression with a segregation free energy of -97.8 kj mol"1 and a saturation grain boundary concentration of 28.0 ng cm"2.

Conclusions

It has been shown by the fact that the thermodynamic parameters derived from our measurements of sulfur grain boundary segregation in nickel are in very good agreement with other studies from the literature that accurate quantitative analysis of equilibrium grain boundary segregation can be achieved using SIMS on bulk samples. The main advantages of NanoSIMS for the study of grain boundary segregation can be summarised as follows:

1. The limit of detection is as low as a few tenths of percent of a monolayer 2. The sample preparation is very simple - conventional mechanical polishing 3. No grain boundary fracture is required in contrast to Auger Electron Spectroscopy. 4. Any type of interface can be analyzed (general or special grain boundaries, twin

boundaries).

Acknowledgements

KLM is grateful for support from UK EPSRC platform grant EP/F048009.

References

[1] A. Larère, M. Guttmann, P. Dumoulin, C. Roques-Carmes, Ada Metall. (1982) 30 , 685.

[2] L. Ben Mostepha, G. Saindrenan, N. Barbouth, A.M. Brass, J. Chêne, (1990) Scripta Metall. Mater. 24 ,112:.

[3] M.P. Seah, Auger and X-ray Photoelectron Spectroscopy, vol. 1, Wiley, New York, 1990.

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[4] L. Karlsson, H. Norden and H. Odelius, Acta Metall. (1988) 36 , 1.

[5] K. L. Gavrilov, S. J. Bennison, K. R. Mikeska, R. Levi-Setti, Ada Mater. (1999) 47 , 4031.

[6] N. S. Mclntyre, C. M. Huctwith, K. F. Taylor, E. Keating, N. O. Petersen and A. M. Brennenstiihl, Surf. Interface Anal. (2002) 33 , 447.

[7] S. Lozano-Perez, M. Schroder, T. Yamada, T. Terachi, C.A. English, C.R.M. Grovenor, Appl. Surf. Sei. (2008) 255, 1541.

[8] N. Valle, J. Drillet, A. Pic, H.N. Migeon, Surf. Interface Anal. (2011 ) 43, 573.

[9] D. McLean, Grain Boundaries in Metals, Clarendon Press, Oxford, 1957.

[10] A.B. Vladimirov, V.N. Kaigorodov, S.M. Klotzman, l.S. Trakhtenberg, Fizika Metallov i Metallovedenie ( 1975) 39 , 319.

[11] F. Christien, C. Downing, K. Moore and, C.R.M. Grovenor (2011) submitted to Surf. Interface Anal.

[12] P. Lejcek, A. Rar, S. Hofmann, Surf. Interface Anal. (2002) 34 , 375.

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Supplemental Proceedings: Volume 1 : Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

The periodic unit of doubly-diffracted reflections from periodic grain boundaries in cubic crystals and its relationship with coincident site lattice

M Shamsuzzoha School of Mines and Energy Development, The University of Alabama, Tuscaloosa, AL 3548

Keywords: Double diffracted unit of reflections, periodic grain boundaries, CSL

Abstract

A geometrical formulation of how the double diffraction scattering from constituting crystals of a periodic grain boundary in cubic crystal yields a doubly-diffracted reflection unit within the composite diffraction pattern of the boundary is provided. The doubly-diffracted reflection unit formed represents a primitive displacement shift complete unit of zero-layer reciprocal lattices (DSCZLRL) of the coincident site unit of zero-layer reciprocal lattices (CSZLRL). Its reciprocal transformation yields a real space unit of coincident site lattices that includes a coincident site of zero-layer lattice (CSZLL) of the periodic boundary as its basic unit. The formulation is applied in a case study on the determination of the coincident site lattice (CSL) from the double diffraction electron pattern of a 13 [101]/(-111) grain boundary of aluminium.

Introduction

Electron diffraction study of a two- or multi-phase material often encounters diffraction spots originating from a phenomenon known as double diffraction. These double diffraction spots usually exist as spurious reflections. A composite electron diffraction pattern containing such double diffraction spurious reflections brings difficulties in structural analysis. To overcome such difficulties a systematic investigation of the diffraction pattern to identify the real diffraction spots from those of double diffracted diffraction spots is required. However, double diffraction spots are not always undesirable, and in some instances they provide useful information, such as the production of "moire" fringes in the transmission electron microscopy image. Double diffraction spots have been successfully used in pioneering work on the determination of the unit cell of ordered compounds such as AuCu(lI) [1, 2] and Au3Mn [3]. More recently analyses of double diffraction spots taken on epitaxial a-Fe2Ü3 islands grown on (0001) Al203[7] have conclusively determined the side of the substrate at which the epitaxial feature lies.

For diffraction of an electron beam from a crystal, double diffraction spots arise in a process in which an initially diffracted beam passing through the subjective crystal assumes the role of an incident beam and produces more diffracted beams. Accordingly, an initially diffracted beam with diffraction vector ri from a crystal can be re-diffracted by the crystal to yield a diffracted beam with diffraction vector r2. In assuming the role of the incident beam the diffraction vector rr causes a change of origin and results in a doubly-diffracted beam at a diffraction vector n-r2. A zero-order Laue zone (ZOLZ) diffraction pattern is the zero-order reciprocal lattices (RL) of the zone where each diffracted beam represents a reciprocal lattice point. The basic unit for a zero-layer RL pattern can be described in terms of two non-parallel basis vectors. Hence, a difference between position vectors of any two RL points in zero-layer reciprocal lattices can be

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obtained by linear combination of the two basis vectors of the unit. This implies that a difference in position vector of two RLs in a zero-layer pattern is also a RL vector. The RL representation of a diffraction pattern as discussed earlier indicates that the diffraction vector n-r2 associated with a doubly-diffracted beam represents an existing diffracted beam. This implies that a single crystal containing a periodic unit of lattices does not yield double-diffracted beams in its reciprocal space and therefore cannot produce additional RL points in the reciprocal space except at those that are not permitted by the extinction conditions of the space group of the crystal.

In metals and alloys, crystals of the same species of random mutual orientations form what is generally termed as "general grain boundaries". In such boundaries, lattices of opposite species existing across the interface are not known to form a periodic unit. The double diffraction phenomenon in such materials usually yields doubly-diffracted beams as additional reflections and creates complicated diffraction patterns. However, periodic boundaries in such materials yield periodic unit of lattices and RLs of the participating crystals. Double diffraction patterns of such boundaries, owing to their formation, are likely to exhibit certain periodicity. This paper provides two-dimensional analyses of how doubly-diffracted reflections from a periodic boundary assume a periodicity that can be utilized to obtain information about their related coincident site lattices (CSL). The paper also provides a case study on the determination of the CSL from the double diffraction patterns of a 13 [101]/ (-111) grain boundary in an Al sample.

Origination of Double Diffracted Reflections from CSZLRL of a Periodic Boundary of Cubic Crystals

In electron diffraction scattering of a periodic boundary, development and distribution of doubly-diffracted reflections are strongly dependent on the position of fundamental diffraction spots arising from the constituting grains of the boundary. Owing to a periodic arrangement of lattices across the periodic boundary diffraction spots of the constituent crystals in the resulting diffraction pattern assume a common super-unit. A zero-order Laue zone (ZOLZ) diffraction pattern represents the zero-layer reciprocal lattices (ZLRL), even though the extinction conditions of reflection due to the Bravais lattice of the real space crystal may prohibit the presence of many zero-layer lattices. Hence, the zero-layer lattice representing a composite diffraction pattern of a periodic boundary exhibits a coincident site of zero-layer reciprocal lattices (CSZLRL) unit. Geometrical parameters of such a CSZLRL unit of a periodic boundary can be utilized to locate the position of doubly-diffracted reflections in reference to the fundamental reflections of the constituting crystals of the boundary. In this respect, identical crystal symmetry of real and reciprocal lattices of cubic crystal enables a recently developed method [5] of constructing CSLs for cubic crystals to be used for geometrical construction of such CSZLRLs. Following this method [5], the CSZLRL unit can be designated by an odd integer known as Sigma (S), which is the magnitude of a reciprocal lattice vector termed as the "sigma generating vector" (SGV) present in the ZLRLs. The ZLRL for an arbitrary orientation of a crystal assumes a network of orthogonal (square or rectangular) units of reciprocal lattices. The unit can be either primitive or centered. The SGV (Gi) in such a ZLRL network of primitive orthogonal unit yielding the magnitude of an £ can be expressed in dimensionless form as:

G, [S mT] ! (1)

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where S and T are unique non-prime integers, m is the ratio of the length of the cell edges of orthogonal RL unit, and i and j are the unit vectors along the cell edges of the orthogonal unit. For a centered orthogonal RL unit, obtaining such a position vector by equation 1 where S ± T is odd requires dividing each cell edge of the orthogonal unit of the participating crystals by 2. No division of such cell edges is needed for the case where S±T is even. Hence the magnitude of the SGV remains identical irrespective of whether the orthogonal network of the ZLRLs is primitive or centered. Such a determination of the SGV for a network of square RL units for a [001] orientation of a cubic crystal is illustrated in Figure la. In the square network of RLs where m=l the magnitude of the SGV assume an odd integer, 5 when S=3, T=4 . The SGV (Gr) thus determined is shown in the figure as OA.

The cell edges of the CSZLRL unit obtained by this method are originated from another vector, which is the square root vector of SGV (Gi ). The square root vector of SGV is also the summation of two other vectors each of which yields different or identical unique non-prime integer. In dimensionless form the square root vector\/( Gi) of SGV and its normal coplanar vector ( G2) forming the cell edges of the ZLCSRL unit can be expressed as

757 = [S' mt]

fm'T -mS'

(2)

(3)

Figure 1. (a) Schematic representation of SGV and associated sub-lattice unit formed by the square root vector of SGV for a zero-layer reciprocal lattice of primitive or body-centered cubic crystal, as viewed along [001]. (b) Schematic of a I 5 CSZLRL that is formed due to a mutual rotation around a common [001] of two primitive or body-centered cubic crystals of the same species. where S' and V are co-prime integers that may or may not

be equal. Equations 2 and 3 are also valid for the network of centered orthogonal lattice units. In these cases a determination of the square root of either position vector^ Gi) or (VGi) with S' ±T being an odd integer, has to be made on the basis that the magnitude of each participating cell edge is one half of that for the unit. Translational symmetry of the RLs also allows the pair of vectorsx/ Gi and VG2 to define a sub-lattice unit designated as OBCD in Figure la. Geometrical analysis of thus constructed unit originated from the primitive orthogonal network of ZLRLs reveals the following relationships between VGi, VG2 and m:

| V G , | / | V G 2 | = m (4)

A CSZLRL unit constructed from such a unit of VGi and VG 2 for the square RL unit of Figure la by the operation described in the method [5] is shown as OPQR in Figure lb.

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Displacement Shift Complete of Reciprocal Lattice (DSCZLRL) Unit of Periodic Boundary and Associated Doubly-Diffracted Unit of Reflections in Related Diffraction Pattern

The CSZLRL described above has its coincident and non-coincident RL sites forming another type of partly-filled network of ZLRLs by a geometrical method described by Balluffi [6]. The unit is termed as a displacement shift complete of reciprocal lattices (DSCZRL). Such a DSCZLRL unit within the CSZLRL unit of a [001] oriented 1=5 (210)/ [001] periodic boundary given in Figure lb is designated as OPQR. Geometrical analysis of presently constructed CSZLRLs reveals that the axial vectors (Di and D2) of DSCZRL are related to the cell parameters (I; and lj) of the unit of basic RL network via the sigma generating vector (I ) given in equation 1. For a primitive orthogonal network of ZLRLs the relationships are:

Di = li/Vl ; D2 = lj/Vl (5)

For a centered orthogonal network of RLs the relationships assume the following forms:

Di = li /Vl ; D2 = lj /VZ if S' ± T' is an even integer (6a) Di = h /2Vl ; D2 = lj /2Vl ifS '±T' is an odd integer (6b)

A similar geometrical analysis also yields relationships between axial vectors (̂ iGi and VG2) of ZLCSRL unit and the axial vectors (li and lj ) of the primitive orthogonal network of the ZLRL unit as:

VGI = Vl li ; L2 = VG2 lj S' ± V is an even integer (7)

For a centered orthogonal network of RLs the relationships become:

VGi = Vlli ; VG2L-Vllj S'±T' is an even integer (8a) VGiL = 2Vlli ; VG2 = 2Vsij S' ±T' is an odd integer (8b)

Combining (5) to (8) provides the relationships between the axial vectors for DSCZRL and the CSZLRLs in terms of I as:

Di= VGi Z ; D2= VG2S (9)

The relationships expressed in (9) reveal that the axial vector of DSCZLRL always intercepts the respective axial vectors by I irrespective of whether the orthogonal network of ZLRLs is primitive or centered. It also reveals the existence of parallel axes relationships between DSCZLRL unit and its related CSZRL. This unit, although strictly not a primitive unit (because each cell edge of the unit is not occupied by a RL), can be considered as another basic unit of the CSZLRL unit of the boundary

For a CSZLRL, ei* and ej* can be the basis vectors that have the magnitude of Di and D2 for the network of thus constructed DSCZLRL unit for a specific X value (which assumes 5 for the CSZLRL unit of Figure 2). A position vector (r*) representing a point on an arbitrary node of this DSCZLRL lattice network in the CSZLRL can be expressed in terms of ei* and ej* as:

276

r* = nei*+Çej*(ll) (10)

where r\ and £ are non-prime integers that assume values ranging from 0 to S. Accordingly, a difference in position vector of any two RLs i.e. (r,* - r2*) of the DSCZLRL network can be expressed as:

r 1 * - r 2 * - ( V r | 2 ) e i + ( ^ - y e j (11)

In view of each of n and I being a non-prime integer with a F i8u r e ,f . Doubly-diffracted 1 .i , r u i. n . Twt. <. • «.v. .u e reciprocal lattice points shown as

value that lies between 0 to I , the terms m the parentheses of snM£, fi„ed d r c £ s forming (he equation 11 also becomes a non-prime integer with a value primitive DSCZLRL units within that also lies between 0 to S. and can be expressed as: a CSZLRL unit of a

[001]I5 (210) grain boundary in v n 2 = i ; ÇrÇ2=Ç <12> asimP,ecubiccrystal-

Combining (10), (11) and (12) reveals that the difference in position vectors (n* - r2*) of two arbitrary nodes of a DSCZLRL is a position vector (r*) of a different node of the DSCZLRL lattice network. This indicates that a linear combination of axial vectors ej* and ej* of a DSCZLRL unit produces RL point in all empty nodes of the DSCZLRL. The process thus makes every node of the DSCZLRL to fill by a RL. In other words, the process of a linear combination of axial vectors t-,* and ej* just discussed allows the CSZLRLs to transform into a primitive basic unit of coincident RLs, i.e. primitive DSCZLRL. Thus developed primitive DSCZLRL unit marked as OEFG in the CSZLRL of the [001] oriented L=5 (210)/ [001] grain boundary is shown in Figure 2.

Doubly-diffracted reflections present in a diffraction pattern are additional to those of the fundamental spots that arise due to the initial diffraction of constituting crystals. These additional diffraction spots do not take part in formulating a coincident unit formed by fundamental diffraction spots. However, each of such additional reflections represents the difference vector of two diffraction vectors. In reciprocal space such a difference in two diffraction vectors is represented by a position vector that is also a difference between two RL vectors. Occurrences of this type of difference in RL vectors are shown in earlier discussion to result in the formation of a network of a primitive DSCZLRL unit. The application this operation within a diffraction pattern of a periodic boundary also brings the difference between diffraction vectors for all diffracted beams to resulting in the formation of a network of a basic primitive unit (that is analogous to that of the primitive DSCZLRL unit) discussed earlier. Each reflection of the unit then can be considered to be arising from both crystals that lie across the interface. This unit of reflection can therefore be termed as coincident unit of doubly-diffracted reflections (CUDDR) unit. In this respect, the primitive DSCZLRL unit (OEFG) of Figure 2 can also represent such a CUDDR unit of the doubly-diffracted reflections present in the diffraction pattern of a [001] oriented S=5 [001]/ (210) grain boundary.

277

Co-Relation of CUDDR Unit of Diffraction Pattern and the CSZLL

The vectors ei*, ej* discussed above also become the basis of the primitive DSCZLRL. These basis vectors are common to the lattice of both constituting lattices for the periodic boundary. Hence, the reciprocal transformation of these basis vectors results in a super-lattice of coincident lattices in real space. This also implies that the CUDDRL unit of a composite diffraction of a periodic boundary upon reciprocal transformation results in a super unit of coincidence lattices. Thus reciprocally-transformed real space super unit of coincident lattices can be either the actual coincident unit or an integral multiple of the actual coincident unit. Further information about this transformed real space lattices can be obtained from the relationships between ej*, ej* and their respective VGI*, \G2* of the CSZLRL. The relationships of these d*, ej* for the primitive DSCZLRL unit with their corresponding axial vectors VGI*, VG2* of the primitive CSZLRL can be obtained by equation 9 as

ei* = VGi*/L ; ej* = VG2*/£ (13)

Equation 13 reveals that the reciprocal transformation of ei* and ej* results in a virtual super-lattice unit of the zero layer coincident lattices in real space This unit also includes another real space lattice unit that is formed due reciprocal transformation of CSZLRL of RLs. The latter real space lattice unit is a primitive DSCZLL of the virtual super lattice unit of the zero layer lattices.. This indicates that the reciprocal transformation of a CUDDR unit in a composite diffraction pattern yields a zero-layer virtual super-lattice, but the same transformation of the unit formed by fundamental reflections brings a zero-layer sub-unit of lattices as the basic unit of the vtrtual super-lattice unit However, the basic unit is formed by the reciprocal transformation of the unit formed by those fundamental reflections that are present after satisfying the conditions for extinction of reflection due to the Bravais lattice of the crystal. This allows the basic unit to be the smallest zero-layer lattice unit, i.e. coincident site of zero-layer lattice (CSZLL), if no conditions for extinction exists in the diffraction pattern due to the Bravais lattice. On the other hand, existence of conditions for extinction of reflections in the diffraction pattern due to Bravais lattice allows the basic unit to divide one or both cell parameters by an integer (derived from the extinction conditions) to yield related cell parameter or cell parameters of the CSZLL. In any of such cases just described, each cell edge of zero-layer virtual super-lattice forms by reciprocal transformation of CUDDR always be the integral multiple of respective cell edge of thus determined CSZL The information provided above for the virtual super lattice unit along with other crystallographic data such as the presence of rectangular CSZLRL units in the diffraction pattern as given in Table 1 can be used to the determine the real CSZLL of periodic boundaries.

Determination of the CSL from a CUDDR Unit of the Experimental Double Diffraction Pattern of a £3 [101]/ (-111) Grain Boundary of Aluminium

A nano-diffraction taken from a S3 [101]/ (-111) symmetrically tilt grain boundary in aluminium is shown in Figure 3a. Its schematic is also shown in 3b. The pattern exhibits both fundamental and double-diffracted spots. Fundamental diffraction spots due to bi-crystal of the boundary are indexed. The unit of fundamental diffraction (representing the CSZLRL unit) and the CUDDR unit present at the origin of the pattern are outlined and designated as OHIJ and OKLM,

278

respectively. The axial vector of CUDDR (OKLM) intercepts the respective axial vectors of the CSZLRL (OHIJ) by S = 3. Analysis involving reciprocal transformation of the CUDDR yields a zero-layer super-lattice formed by -333, -6 6 -12 and their symmetry-related lattice points and is outlined as OSTU in Figure 3c. A similar analysis of the OKLM unit results in another coincident lattice unit that is bounded by lattice points -111, -22-4 and their symmetry related lattice points. Table 1. Data in relation to the construction of CSZLRLs for cubic crystals o

1

3

5

7

9

11

13

Rotation axis

001

no

001

11]

no

110

100

Type1

Primitive square (in BCC) a*, a* Centered square (in FCC)

a./V2,a>/V2 Primitive rectangle (in FCC)

a*,a*/V2 Centered rectangle (in BCC)

a.a-^2

Primitive square (in BCC) *o. a*

Centered square (in FCC)

a«5J. <"/V2

Centered rectangle (inBCC)a>V2,a«V6

(in FCC) a - ^ j , a«V5/2

Primitive rectangle (in FCC) a. ,a»/V2

Centered rectangle (in BCC) a, o«V2

Primitive rectangle (in FCC) a,« a*V2

Centered rectangle (in BCC) a * ,a«V2

Primitive square (in BCC) *a, a*

Centered square in (FCC)

m

1

1

2 -1 2

1

1

3

V 3

2

2

2

2

1

1

S

2

2

1

1

3

3

1

1

3

3

7

7

12

12

T

0

0

2

2

4

4

4

4

6

6

6

6

5

5

S'

1

1

1

1

2

2

2

2

1

1

3

3

2

2

T'

2

2

1

1

3

3

O'O

0

70 5

70.5

36.9

36.9

81 8

818

70 5

70 5

50.5

50.5

22 6

22.6

Parameters'

No CSZLRL

rWî .Ca .A^lVz:

a . / 2 v X ( a . i / z ) ) A

a»VS,a«VS

$w (a»V2)VT,(a • V6)VS

a . VI, (a ./V2) Vl

a»VS,o-V2VS

a . V £ , a . V 2 i / l

a«V5,a./V2V5

a + VË, a * VË

&W

Direction3

<111>, <112> <112>, <lll> <210>, <210>

<3I0>, <3I0> <123>, <145>

<123>, <145>

<122>. <U4> <114>, <122> <113>, <233> <233>, <113> <230>, <230>

<150>, <I50>

1. Type of two-dimensional reciprocal lattice network unit with the dimension of two cell edges expressed in terms of unit cell length a* of the crystal 2. Lattice parameters of CSZLRL expressed in terms of the a* and their plane. 3. Direction <uvw> of the zero-layer reciprocal lattices in terms of the lattice parameters of the constituting crystals.

It is outlined as OVWX in Figure 3c. The cell edges OM and OK of the related CSZLRL unit (OKLM) yields a OM/OK ratio of 2̂ 2 . According to Table 1, a 1-3 boundary possessing this CSZLRL with a ratio of 2V2 for such cell edges is evolved from a centered rectangular reciprocal lattice unit. In this context it is worth noting that a network of a centered rectangular unit of reciprocal lattice can only arise from a reciprocal unit that has a body-centered Bravais lattice. The respective cell edges (OX and OV) of the real space lattice unit (OVWX), which are reciprocally transformed from the CSZLRL (OKLM) unit, also yield ratio of 2\/2. The 2V2 for the ratio indicates that the reciprocally transformed real space coincident unit OVWX also arises from the lattice network of a centered rectangular unit, which evolves from a cubic crystal that has a body centered Bravais lattice. This implies that the reciprocal transformation of a BCC crystal yields a BCC crystal; but this is contrary, because a CSZLRL unit made of crystals of a BCC Bravais lattice is expected to yield a coincident unit of lattices that are formed by the lattices found in a FCC crystal [5]. Table 1 reveals that the CSZLL formed by the lattices of FFC

279

crystal yields a ratio of V2 for the division of such cell edges. This indicates that the ratio of such cell edges of the real CSZLL for the experimental 13 [101]/ (-111) symmetrically tilt grain boundary should be V2 instead of 2V2 [5], Taking allowances of this modification in the ratio of the cell edges of the lattice unit of OVWX, the actual CSZLL that lies within the super unit of OSTU of Figure 3c has to be bounded by -111, -112 and their symmetry-related real space lattices. Thus determined unit of coincident lattices for the boundary is outlined as OV'YZ in Figure 3 c. This lattice unit of coincident lattices is the basic unit that can be present in the super-lattice unit of OSTU and is the CSZLL. The addition of a third axis along the rotation axis i.e. [101] of the CSZLL yields a three dimensional CSL.

Figure 3. (a) Zero-order Laue zone diffraction pattern of the [101]I3 (-111) boundary taken with electron beam parallel to [101]. (b) Schematic of (a); small circles are due to double diffracted rays, which form the primitive DSCZLRL unit of the CSZLRL. (c) Schematic of relationships between the CSZLL (OVYZ) and super-lattice unit arising from reciprocal transformation of the DDRL unit of the experimental diffraction pattern in (a).

Conclusions

Analysis on double diffraction scattering of the constituting crystals of periodic boundaries re-veals that a primitive unit of doubly-diffracted reflections becomes the basic unit of reflections within the composite diffraction pattern of the boundary. The presence of such a unit in the expe-rimental diffraction pattern of a boundary indicates the existence of lattice periodicity across the interface. Since the basic unit of doubly-diffracted reflections acts as the primitive DSCZLRL unit of the CSZLRL unit formed by fundamental reflections, the reciprocal transformation of the unit yields a virtual super-lattice unit of real lattice in which the reciprocally transformed unit due to fundamental reflections assumes the role of the basic unit of real lattices. Attainment of this virtual super-lattice unit from the doubly-diffracted unit provides a way to accurately estab-lish the CSZLL of the boundary.

References

[1] A. B. Glossop and D. W. Pashely, Proc. Roy. Soc. A 250 (1959), 132. [2] S. Ogawa, D. Watanabe, H. Watanabe and T. KomodMcta. Cryst. 11 (1958), 872. [3] D. Watanabe, J. Phys. Soc. Japan 13 (1958), 535. [4] A. Tiertz, C. B. Carter and S. McKarnen, Ultramicroscopy 60 (1999), 241. [5] M. Shamsuzzoha. Supplementary Proceedings: Volume: 3: General Paper Selections, IMS (The Minerals, Metals, and Materials Society) 2011, 463. [6] R. W. Ballufïi. Interfacial Segregation, A.S.M., Metals Park, Ohio, (1977)

280

Supplemental Proceedings: Volume 1 : Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

The in-situ Intrinsic Stress Measurements of Cu and Al Thin Films

Jun Young Yu, Youngman Kim

Dept. of Materials Science and Engineering Chonnam National University, Gwangju, 500-757, Korea

Keywords: thin film, in-situ, intrinsic stress

Abstract

We observed the in-situ stress evolution of Cu and Al thin films during deposition on (111) Si

wafers using a thermal evaporation method in terms of deposition rates. Cu and Al films were

deposited at rates ranging from 0.5 to 1.5 Â for Cu and for Al. In-situ stress values in the

growing films were obtained using multi-beam curvature measurement system installed in the

chamber of thermal evaporator. For the copper films the in-situ intrinsic stress showed a typical

three step behavior of initial compressive, tensile and gradual compressive stresses. For the

aruminuim films the in-situ intrinsic stress showed rather unstable behavior without highly

reliable repeatability.

Introduction

In general, the internal stresses in thin films in the presence of electronic components used in the

film due to the crack and cause an electrical short causing adverse effects on component life, and

the band-gap semiconductor material increases the resistance to move (shift) to happen is also.

Cutting tool wear is also applied to the case of thin-film adhesion and wear resistance by a

unique stress reduces degradation of mechanical properties, such as are taken. Films with

increasing film thickness tends to increase stress, and this will limit the thickness of thin film

deposition, because. In addition, deformation of the film by the substrate can lead to stress, but in

most cases relative to the substrate film really does not matter, because many so thick. However,

in integrated circuit technology, silicon wafers, finely curled a precise tolerances in a lithography

processis very difficult to maintain. And to continue shrinking the size of the device because of a

very thin film thickness in the area is very important to control stress is being recognized. [1]

Film stress is classified as thermal stress and intrinsic stress. The thermal stress between the

substrate and the film is caused by differences in thermal expansion coefficient. The unique

281

stresses that occur during the growth of the film stress is defined as the growth referred to as

stress. This thermal effect appears independent of stress occurs, the current mechanism has been

proposed that various stresses, intrinsic stress of exercise, depending on the source material

difference in their behavior seems to exercise with high-mobility of Volmer-Weber type growth

in copper, silver, gold, and metal film thickness increases, the film based on the 'initial

compressive stress ', 'tensile stress', 'second compressive stress' levels of stress, its unique three-

phase behavior has been shown [1,2]. Step-by-step behavior of thin films of these stresses on the

initial compressive stress, Cammarata argues source material be deposited onto the substrate, the

islands are created early in the process created a new surface that caused the growth of thin films

due to surface effects of compressive stress receive initial islands , and since the island is firmly

attached to the substrate and the equilibrium lattice constant is no longer able to keep the films

showed compressive stress. [3,4] Spaepen argues due to manufacturing thin film to stable atomic

position rather than incursive ad-atom has amount that the initial compressive stress occurs [5].

Tensile stress of the film, the Hoffman when separated after the merge of the surface energy is

greater than the energy of grain boundaries resulting in the film being an island, merging by

volume reduction takes place when there argued that the tensile stress. [6] These argues Nix and

Clemans, and improved by Seel and Thompson, through simulation studies are underway. [7-21]

Case of Second compressive stress, Spaepen that the island growth during thin film deposition

layer is compressed to meet two extra atoms in the process created sufficient space to enter the

normal state than in the film is more compressive stress causes the atoms emerges argued that.

[22] Chason a continuous thin film layer deposited films at grain boundaries because of the non-

equilibrium surface chemical potential is added to the grain boundaries of polycrystalline thin

film of atoms penetrates argued that while the compressive stress occurs. After stopping the

deposition of thin films for tensile stress occurs in the thin film deposition flux of atoms to reach

the floor by a reduction of the potential changes in surface chemistry, and hence the atoms

diffuse to the grain boundaries to move back to the surface occurs, contributing to the tensile

stress argues that. [23] Such a mechanism would suggest that the stress at each step many

experimental studies [7-21], despite being a compressive stress generated is still not understand

about the causes is lacking.

In this study, the deposition process occurring in film compressive stress mechanisms in order to

understand in-situ multi-optical curvature measurement system, the deposition rate changes and

282

the deposition of a substance to exercise the film's stress on the behavior how to effect learn

about it analysis of the copper and aluminum films on a unique instrument to study the stress is

raised.

Experiment Procedures

Devices used in this experiment, as shown in Figure 1 are part of larger thin-film and thin film

stress measurement has two parts.

Figure. 1. Schematic diagrams of (a) the thermal evaporation chamber and (b) the in-situ multi-

beam curvature measurement system.

Thin film deposition using the thermal vapor deposition of copper and silver films were

deposited, deposition rate of copper and aluminum films to learn the impact of stress on the

behavior of thin film deposition rates, respectively 0.5 -1.5 Â/s were deposited by varying with.

Multi-optical curvature measurement system while thin-film devices in using the curvature

measurement film thickness changes were observed according to the change of curvature, the

thickness of the film changes Sycon Instruments, STM-100/MF using observed real time during

deposition. Stoney equation [25], measured using thin-film stress was calculated from the change

of curvature.

B.tt

Wf (1)

Here, cav mean stress of the film, Es biaxial elastic modulus of the substrate, ts of the substrate

thickness, tf the film thickness and k is the change in curvature of the substrate. Substrate a 4-

283

inch Si (111) wafers were used. Sensitive to the radius of curvature measured to be 100 urn thick

wafers were polished on both sides, 5 minutes each in acetone and alcohol, distilled water,

ultrasonic cleaning for 10 minutes, and the specimen was dried with nitrogen gas. The pressure

inside the reactor of less than 5 x 10-6torr below maintain high vacuum, and all experiments

were in progress at the same pressure. The temperature was maintained at room temperature

inside the reactor substrate and the deposited material, containing approximately 32 cm, the

distance between the W boat always was kept constant.

Results and Discussion

1. Cu thin films stress behavior in deposition rate change

1.1 Cu thin films initial compressive stress

In this study, thin film deposition rates the impact of stress on the behavior of copper and

aluminum films at 0.5 ~ 1.5Â/s to the deposition rate by changing the behavior of the film stress

changes were observed. Compressive stress in thin films of copper and aluminum causes that

contribute to demonstrating the effect of deposition rate to the maximum value of compressive

stress in the thin film thickness and this thickness, the average stress value was measured. The

maximum value of compressive stress, then, to find the thickness of the non-linear regression

analysis were used. Results for copper films in Figure 2, and 3. In this experiment, copper and

aluminum films, respectively 0.5 ~ 1.5À/S deposition rate was deposited. Figure 2. (a), each

point with a constant deposition rate of deposition is a deposition process. Of copper thin films,

the maximum compressive stress, the deposition rate of change in value almost did not show the

effect of increasing. And Figure 2 (b) at a deposition rate as a maximum value of compressive

stress in the specimen appearing in the film thickness difference can be seen that a nearly

invisible. Therefore, thin film deposition rate of copper thin films causes compressive stresses,

but to contribute to the key variables can not be said.

1.2 Cu thin films incremental compressive stress

Figure 3. in the 45nm film thickness due to the deposition rate of copper films in the incremental

compressive stress shown the value of. Figure 3 for each point shown in the copper thin film 0.1

~ 1.5 Â/s within the scope of the specimens deposited in a uniform deposition rate shows the

values. The value of each point due to the increase of film thickness Of * tf a graph showing the

value sought by the linear regression analysis. Effect on the compressive stress in order to

minimize the impact of the film thickness of copper film thickness in the range of 40 ~ 50nm

284

linear regression analysis was performed. Thin film deposition rate of copper increasing

incremental compressive stress values tended to decrease.

Figure 2. (a) The average stress at the compressive maximum versus deposition rate for copper

thin films, (b) Film thickness at the compressive maximum versus deposition rate for copper thin

films.

Figure 3. The incremental stress at the thickness of 40~50nm versus deposition rate for copper

thin films.

2.A1 thin films stress behavior in deposition rate change

285

2.1 Al thin films initial compressive stress

Figure 4, 5. for each point that appears in the aluminum film 0.1 ~ 1.5 A/s within the scope of

the specimens deposited in a uniform deposition rate shows the values. Aluminum thin films, the

maximum compressive stress, the effect of increasing the value of the deposition rate almost did

not. Figure 4. (a), depending on the deposition rate as shown in the extract was virtually no

change in maximum stress. But Figure 4. (b) increase the deposition rate is maximum, depending

on the value of the initial compressive stress tends to increase when the thickness was slightly.

This is seen as an aluminum thin film deposition rate of the size of the compressive stress does

not contribute to the maximum compressive stress is expected to slow down when it appears.

Therefore, an aluminum thin film deposition rate on the size of the compressive stress does not

contribute to determine when to generate a compressive stress can be thought of as one of the

variables.

Figure 4. (a) The average stress at the compressive maximum versus deposition rate for

aluminuim thin films, (b) thickness at the compressive maximum versus deposition rate for

aluminuim thin films.

2.2 Al thin films incremental compressive stress

Figure 5. in the 30nm film thickness due to the deposition rate of copper films in the secondary

compression stress (incremental compressive stress) is shown the value of Figure 5. for each

point that appears in the aluminum film 0.1 ~ 1.5 A/s within the scope of the specimens

deposited in a uniform deposition rate shows the values. The value of each point due to the

286

increase of film thickness of tf a graph showing the value sought by the linear regression

analysis. Effect of compressive stress in the film thickness in order to minimize the impact of the

aluminum film thickness in the range of 20 ~ 30nm linear regression analysis was performed.

Aluminum thin film deposition rate increases in the incremental compressive stress values

appear to have a significant impact.

Figure 5. The incremental stress at the thickness of 20~30nm versus deposition rate for copper

thin films.

Conclusion

Using thermal evaporation deposition of copper and aluminum films on a silicon wafer multi-

beam curvature measurement system growth in real time using the device due to the increase of

film thickness were observed stress behavior. Both copper and aluminum thin film thickness

increases, the initial compressive stress and tensile stress, incremental compressive stress,

mobility of the unique three-stage manifestation was observed. In addition, the deposition rate of

the impact of stress on the behavior of each film deposition rate to Read 0.1 - l.SÂ/s with

varying stress behavior was observed accordingly. Both copper and aluminum thin film

thickness or the maximum compressive stress, maximum compressive stress values appear as the

relationship between the deposition rate did not show a particular trend. However, the deposition

rate of copper films with increasing incremental compressive stress the value of the lower tended,

aluminum thin film deposition rate increases, the incremental compressive stress appears, had a

tendency to increase in thickness.

287

Refernces

[I] Ryu Sang, Lee Kyungchun, Kim Youngman, Kor. J. Met, Mater. 46, 5,283 (2008) [2] R. Abermann, R. Kramer and J. Mser, Thin Solid Films 52, 215 (1978). [3] A. L. Shull and F. Spaepen, J. Appl. Phys. 80, 6243 (1996). [4] M. Laugier, Vacuum 31, 155 (1981). [5] R. C. Cammarata, T. M. Trimble and D. J. Srolovitz, J. Mater. Res. 15, 246 (2000). [6] E. Chason, B. W. Sheldon, L. B. Freund, J. A. Floro and S. J. Hearne, Phys. Rev. Lett. 88, 156103 (2003). [7] R. Abermann, R. Koch and R. Kramer, Thin Solid Films 58, 365 (1979). [8] R. Koch, H. Leonhard and R. Abermann, Thin Solid Films 89, 117 (1982). [9] H. P. Martinz and R. Abermann, Thin Solid Films 89, 133 (1982). [10] R. Abermann and R. Koch, Thin Solid Flms 142, 65 (1986). [II] R. Abermann, Thin Solid Films 186,233 (1990). [12] G. Thurner and R. Avermann, Thin Solid Films 192, 277 (1990). [13] D. Winau, R. Koch and K. H. Rieder, Appl. Phys. Lett. 59, 1072 (1991). [14] S. Ryu, K. Lee, S. Ma and Y. Kim, J. Nanosci. and Nanotech. 7, 4081 (2007) [15] C. Friesen, S. C. Seel and C. V. Thompson, J. Appl. Phys. 95, 1011 (2004). [16] C. Friesen and C. V. Thompson, Phys. Rev. Lett. 89, 126103 (2002). [17] H. Windischmann, CRC Crit. Rev. Solid State Mater. Sei. 17, 547 (1992). [18] M. F. Doerner and W. D. Nix, CRC Crit. Rev. Solid State Mater. Sei. 14, 225 (1988). [19] L. Vecchio and F. Spaepen, J. Appl. Phys. 101, 063518 (2007) [20] C. V. Thompson and J. R, Llyod : MRS Bull. 18 (1993). [21] R.Chaudhari. J. Vac. Sei. Technol., 9, 520 (1972) [22] D. W. Pashley, M. J. Stowell and M. H. Jacobs , Phill. Mag., 10, 127 (1964) [23] F. Spaepen, Acta Mater. 48, 31 (2000). [24] E. Chason, B. W. Sheldon, L. B. Freund, J. A. Floro and S. J. Heame, Phys. Rev. Lett. 88, 156103 (2003). [25] A. L. Vecchio and F. Spaepen, J. Appl. Phys. 101, 063518 (2007)

288

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DELAMINATION CHARACTERIZATION OF BONDED INTERFACE USING SURFACE BASED COHESIVE MODEL

Manivannan Ramamurthi1, Young Suk Kim2

'Graduate School, Kyungpook National University, Daegu, Korea department of Mechanical Engineering, Kyungpook National University, Daegu, Korea

Keywords: Surface based cohesive Model, Delamination of bonded surfaces, 90 degree peel test

Abstract

Element based cohesive zone model (ECZM) is employed widely in studying delamination (or) decohesion of interface surface in adhesively bonded materials using finite element models (FEM). This paper uses surface based cohesive model (SCZM) available in commercial finite element Analysis (FEA) code in the place of ECZM by considering similarities between both models when interface thickness is zero for the advantages of less input parameters, reduced computational time, and easy modeling. Mode I fracture study with 90 degree peel test experiment and simulation are done in polymer coated steel for this analysis. Results are compared with ECZM. Surface based cohesive model predicts delamination well and results are closer to ECZM. Cohesive model requires fine meshing in FEM to predict delamination closely. To verify the feasibility of using it in coarse meshes, same simulations are done with coarse meshes. The results reveal that coarse mesh simulations also predict delamination closer to fine meshes with reasonable accuracy.

Introduction Delamination (or) decohesion in interface of surface bonded materials like polymer coated steel is often occuring problem in industrial applications. If the parameters related to this delamination onset and prorogation (Cohesive parameters) are known, it can be predicted. In other words, delamination characterization is to be done. Cohesive zone model is appropriate for this analysis. Element based cohesive zone model (ECZM) is being widely employed to analyze these kind of delamination problems via finite element analysis (FEA) codes. In recent years, cohesive zone modeling has been implemented in commercial FEA code packages like ABAQUS [1]. Decohesion behavior in interface has been well analyzed in many literatures [2]. There are a number of cohesive zone models available in literature [3]. ABAQUS uses a triangular or bilinear model. Three cohesive parameters are required to define cohesive zone in FEA codes. They are (i) Critical fracture energy Gc required to separate the bonded surface (ii) Maximum nominal stress or cohesive strength Tu|t of the interface and (iii) Stiffness Ketr of the interface. Among these, Gc is the only and most important parameter that should be calculated by experiment [4-8]. Measurement of TUH is difficult and is not well established for most of the cases. However, V. Gupta et al [9] had measured the T„it by laser spallation technique. In FEA analysis, this T„jt value is dependent on mesh size used in FEA. If Tuit is not measureable, penalty approach can be used to find cohesive parameters according to mesh size. T.Diehl [10,11] followed penalty method with only Gc value known. T^t and Keff values were treated as penalty parameters by him. Initial stiffness value Keff is a penalty parameter that should be computed. Many guidelines are available for selection of reasonable Keff value [1,7,10,11]. Improper selection of Keff will lead to numerical problems and or poor stiffness between the surfaces. Thus selection of Keff always needs a reasonable analysis and care. Also, the computation time is high when element based cohesive zone model is used, particularly in 3D models. In ABAQUS version 6.8, a surface based cohesive model (SCZM) capability was introduced. It is applicable where the interface thickness is zero or very less. In this surface based model, keff is not a mandatory input but an optional value. This reduces number of mandatory variables required to define cohesive model into two i.e., Gc and T„it. When interface thickness is zero, the surface model's response is similar to element model by the similarities between these two models. This paper takes this similarity and uses this surface based cohesive model in the place of element based cohesive model to study delamination of zero thickness interface for the advantages of less input parameters, reduced computational time and easy modeling. Delamination between Polyethylene terephthalate (PET) and Polyvinyl chloride (PVC) layers in polymer coated steel in mode I fracture is studied here and related cohesive parameters are derived with only Gc value known. Since Peel test is considered one of the best method to analyze interfacial properties where

289

flexible laminates are involved [12-19], 90 deg peel test is carried out. Cohesive modeling requires fine meshing and cohesive parameters are mesh dependent. In many of the real life problems, fine meshing is not possible. To check the feasibility of applying it in coarse mesh models, the same fine mesh simulations are conducted with two different coarse meshes. Results are compared with element model for cross verification.

Element Based Cohesive Zone Model (ECZM) In Element Based Cohesive Zone Modeling (ECZM), cohesive zone is introduced in FEM as the interface surface to be studied. The cohesive zone being used in ABAQUS is shown in Fig 1. As in Fig 1, cohesive zone is drawn as a curve between traction and separation displacement. In elastic region, an elastic modulus ECfr and stiffness K^e are applicable. In this region no crack initiates. When traction reaches the cohesive strength value Tuit, crack initiates. The displacement related to this point is known as crack initiation displacement 8o. After this, delamination propagates. Fracture occurs at a distance called failure separation distance Sf. After this point, the crack becomes traction free surface. Area under the total cohesive zone curve is called critical fracture energy Gc.

Fig 1. Cohesive zone model. Fig 2. (a) Element (b) surface cohesive model. Fig.3 Peel test configuration.

When we assume 3D problem situation to define cohesive zone in ABAQUS, the Cohesive zone parameters are needed to be given in normal direction(n), first shear direction(s) and second shear direction(t). Critical fracture energy Gc (Gi, Gu, Gui), maximum nominal stress or cohesive strength T (T„,Ts,Tt), initial stiffness K (Kn,Ks,K,) are to be given as inputs. In isotropic condition; values in all directions are same. Thus, Critical fracture energy becomes Gc, cohesive strength is T^ and initial stiffness is IQr in all directions. From Fig 1,GC is connected to 8f and T„it via area of triangle.

G c = ( l / 2 ) T u l t Sf (1) Damage initiation ratio 8^,0 is defined as,

S r a t i o ^ o / S f ( 2 ) Keff is computed from the initial slope as,

Keff=Tui,/5o (3) Effective modulus Eeir is derived from following equation

Eeff=Keffheff (4) The effective thickness hetf of cohesive element is 1 unless otherwise real thickness is known. Gc must be calculated from the experiment. In most of the cases, it is difficult to calculate Tuu by experiments. T„it value is dependent on mesh size in FEM. In Fig 1, As long as the area under the curve Gc is kept constant, the simulations can be done with each value of Tuit by modifying 8f and results can be matched with experiment results. Cohesive parameters related to matched simulation are taken as final cohesive parameters. Normally very fine meshes are required for cohesive simulation and these cohesive parameters are mesh dependent [7,10,11]. Stiffness Keff is a penalty parameter that should also be computed. In ideal condition, the cohesive elements must possess infinite Keff. But, in FEM model, a finite value is to be given. This Kefr value should be high enough to avoid much interpénétration of surfaces and should not be very high to cause numerical problems. Selecting a proper Keff is a tricky work. Next, if the models are run with explicit mode, a mass density must be given as an input for cohesive elements.

Surface Based Cohesive Zone Model (SCZM) In ABAQUS version 6.8, a Surface Based Cohesive Zone model (SCZM) was introduced especially for zero or very less thickness interfaces where thickness effect is not considered for

290

analysis. In this model, the cohesive behavior is given as an interaction property between two surfaces in surface based model. Here, initial elastic stiffness ken-is not a mandatory input but an optional value. Keff value of cohesive surface is assumed automatically by ABAQUS based on the K value of underlying elements. Thus the required number of mandatory inputs to be given becomes just 2 in case of surface based cohesive behavior. Moreover, since the cohesive properties are given as interface properties, this model will not add mass to analysis. So, in ABAQUS explicit mode, when surface based model is used, the mass density need not to be given. Elastic traction seperation, damage initation criteria and damage evolution laws are similar for both surface and element based models. But the traction and separation values are calculated differently. Fig 2 shows surface and element based cohesive model between bulk material 1 and 2. In element based cohesive model, Tuit is equal to nominal stress on cohesive elements, Whereas in surface based model, Tuit is computed as the ratio of contact force to current area at each contact point. When element based cohesive modeling is used, the nominal strain 8 is calculated as follows. In Fig 2(a), the cohesive elements under delamination are shown. When the relative displacement between top and bottom of cohesive layer is taken as d , the nominal strain is the ratio between dto original constitutive thickness t„

e=d/ t 0 (5) As said earlier, when interface thickness is zero, original constitutive thickness to value is taken as 1 in element based cohesive model. In that case, e value is equal to relative displacement between top and bottom of cohesive layer i.e., separation (5). In case of surface based cohesive model, as in Fig 2(b), separation (5) is known as contact separation between surfaces. It is equal to the ratio of relative displacements between slave surface nodes to their respective projection points in the master surface. Thus, incase of zero thickness interface, separation values computed by both the methods are similar. As in Fig 1, area under the curve Gc is common for both models. As separation values are same when interface thickness is zero, 8f must also be same for both element and surface based cohesive behavior. Hence, when Gc and ôf are same for both models, Tuit value also must be same to fulfil the triangle relation in case of zero thickness interface. By considering these similarities, delamination characterization has been done in this study using surface based model in place of element based model for the benifits of less parameter input, less computational time and easy modeling.

Damage Initiation And Evolution Criterion In surface based model, damage initiation is the point where the maximum contact stress Tui, is reached in delamination. It is assigned by maximum stress criterion [1] as follows,

max{((T n) /Tn°), (Ts /T s°) , (Tt /T t ° )}=l (6)

Here, T„ - pure normal direction stress, Ts,Tt - First & second shear direction stresses. Tn0,Ts

0,Tt° represent the highest magnitude of the contact stress when occurred deformation is either purely normal to the interface surface or purely in the first or second shear direction respectively. When the ratio of maximum contact stress reaches the value 1, damage initiates. Tn°,Ts

0,Tt° are to be given as inputs. In isotropic condition, we input Tn° = Ts° = Tt° = T„it. Damage evolution is given by BK (Benzeggagh-Kenane) mixed mode behavior [1]. It has the following form.

G I C + [ ( G I I C - G I C ) { G , W /Gy}11] = G c (7)

Here, Gi , Gn , Gni are the fracture energies in mode I, mode II, and mode III respectively. Gshear=Gn+Giu, G-r=Gshear +Gi. n is a material constant. T| value has been taken as unity thus model response will be insensitive to this parameter. In our case critical fracture energy in all three directions are same i.e., Gc. A scalar variable D will calculate the overall damage process. Condition for D in damage evolution is given below,

[Tn .Otherwise J

T S = ( 1 - D ) T 7 (9)

T t =(l-D)T t (10) Here Tn,T,,T, are the contact stress components predicted by the elastic traction-separation behavior for the current separations without damage. Tn is the stress in pure normal direction, Ts and Tt are the stresses in first and second shear direction respectively. In isotropic condition, all direction

291

stresses are same. A status flag called CSDMG [1] will be assigned to verify whether this condition is satisfied or not in field output. When CSDMG reaches 1, it shows failure. Similarly in element model, this flag is SDEG and all stresses are nominal stresses & strains are calculated from Eqn (5).

Critical Fracture Energy Gc Peel test is considered as one of the appropriate method to calculate adhesive (or) critical fracture energy Gc where flexible laminates are involved. A standard peel test configuration is shown in Fig i. In peel test, two adhesive bonded materials are separated at an angle called peel angle 9 at constant velocity rate. The fixed bottom material is called substrate. The material used to pull to separate the surfaces is called peel arm. From peel test, we obtain an output called peel force, P. From peel test, Gc has been calculated by two methods. One is by classical method and another is by ICPeel program [12,13,14] implemented by Imperial college, London. Classical fracture energy method considers only elastic peeling. ICPeel program is meant for plastically bending and deforming peel test. Reason for calculating Gc by ICPeel program is to ensure large plastic bending and deformation occurrence in peel arm. When a peeling process takes place, the total input energy G is equal to

G=G„+Gn (11) Here, Gc is the critical fracture energy required to separate the bonded surfaces. Gp is the plastic work in bending peel arm. When a peel force P is applied over a width of b at an peel angle of e, the total input energy G can be calculated by,

G=(P /b ) ( l - cos0 ) (12) If only elastic deformation occurs in the peel arm and plastic bending is negligible, then the Geis equal to the total input energy. Thus, incase of elastic peeling, critical fractue energy is computed by classical method with following formula,

G=G c =(P/b) ( l -cose) (13) If peel arm undergoes plastic bending, then the plastic bending work in peel arm Gp should be substracted from total input energy G. In that case, by IC Peel program, the Gc is calculated as,

GC=G-Gp (14) I.Georgiou et aï. [12] have derived the formulation to calculate plastic bending work Gp. The equations required to compute Gp are more in numbers and complex in natrue. It can be referred in Ref [12,13]. Total computation has been coded to an excel program called ICPeel [14] developed by Imperial college, London. It can be used to calculate critical fracture energy easily.

Experimental Analysis In this study, polymer coated steel is analyzed. In this steel, polymer layups have been added over GI (Galvanized Iron) steel substrate. Thickness of GI steel substrate is 0.8 mm. GI steel has been surface treated with epoxy layer of thickness 0.005 mm at the bottom. Above GI steel, first a layer of Polyvinyl chloride (PVC) film of thickness 0.08 mm is coated with a strong adhesive bond of 0.005 mm. Above PVC, a layer of 0.03 mm thick Polyethylene Terephthalate (PET) film is coated using zero thickness aluminum deposition. Delamination occurs between PET and PVC polymer layers. This zero thickness aluminum layer will be defined as cohesive layer. Before peeltest, first tensile test was done for GI steel, PET and PVC. GI steel was tested according to ASTM E8/E8M-09 standard [201. PET, PVC films were tested according to ASTM D882:2009 [21]. Mechanical properties of all above are listed in table I.

Table I. Mechanical properties of GI Steel, PET, PVC Property

Density (p), kg/m3

Young's modulus, E ( GPa) Poisson's ratio Yield strength, o y (MPa)

Yield strain, eY

Tensile strength, oTS (MPa) Total Elongation ( % )

GI Steel 7800 201 0.3

150.1 0.01954

261.2 58.84

PET 1400 3.93 0.3 89.6

0.034 163.6 43.8

PVC 1300 1.30 0.41 25.2

0.049 32.6

262.7

a «

■ • Experiment = 10.16 N

Average plateau of peel force

/

0 3 6 9 12 15 18 Displacement (mm)

Fig 4. 90 Degree peel test experiment result.

90 degree single arm peel test was done according to ISO 8510:1 [22] with peel test machine. Bonded steel and PVC portion was considered as substrate and PET was used as peel arm and was pulled with a steady peel rate of 50 mm/min. As in Fig 4, average peel force calculated

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from experiment is 10.16 N. Calculated Gc value by classical energy method from Eqn (13) is "4±Q.02 kJ/m (standard deviation from experiment) and value of Gc by ICPeel method is 0.1395

I/m . Detailed calculation of Gc by ICPeel method can be referred inri2,13,141. kJ/irr,

Results and discussion FE modeling Fig 5 shows the layout of initial 3D peel test simulation configuration. As there is no deformation in GI steel, it was designed as rigid 3D planar shell. Steel was meshed with R3D4 elements. As there is no delamination and relative movement between steel and PVC, both were tied. PVC and PET were designed as 3D deformable solids. PET was meshed with 3 elements and PVC was meshed with 8 elements in thickness direction. Both PET and PVC were meshed with 8 node three dimensional C3D8R elements. A small unbonded portion of PET was used for bending before peeling. Isotropie condition is assumed in this study. Also as there is no deformation in shear 2 direction in 90 deg peel test, only one mesh element was assigned in width direction for all parts. Doing so reduced computational time drastically. 90 degree peel simulation was performed to predict the peel force exerted by peeling of PET from PVC in polymer coated steel using cohesive behavior in the interface. Same experiment condition was simulated as explained in experimental analysis section. In surface based modeling, between the surface of PET and PVC, surface based cohesive behavior was introduced. The default Kcff selection option was chosen. Thus Keif value was not given us input. In element modeling also, 3D cohesive element layer was introduced as separate layer. Cohesive zone was meshed with 8 node three dimensional COH3D8 cohesive elements. Cohesive zone elements were meshed 5 times denser than bulk elements and they were attached to the surrounding elements by tie constraints.

Table II. Initial FE model configuration (Ref Fig 5) S t r e a m

1 2 3

Common values for Surface and element based cohesive models

(mm)

X

50 50 50

Y

48 47 44

Z

2 3 6

Bulk material mesh element

length lb

0.2 1 2

Cohesive mesh element length

( only for element based cohesive moel) lc ( mm )

0.04 0.2

Fig 5. Peel test simulation initial configuration. "Constants: lb/leratio =5(Only for element based cohesive model)

Peel test simulation with 0.2 mm bulk mesh With respect to Fig 5 and table II, initial configuration of stream No.l was modeled for 0.2 mm mesh simulation lor both surface and element model. First, Simulations were run with Gc by classical energy method i.e., Gc=0.4 ± 0.02 kJ/nr (standard deviation of ± 0.02 kJ/m in peel test). This Gc was varied from 0.38 to 0.42 kJ/m . Since the simulation results were matching closer with 0.38 kJ/m , this value was fixed as final Gc. Simulations were done with varying T„i, by changing Sf. In this study, for all cases, ô> values are taken with respect to bulk element mesh length lib. Values of Tuit for each 8f value were computed by Eqn (1). By analyzing all considerations, simulation was started with 5f= 0.16 mm. For element based model, as discussed, Kefr value was chosen in such a way that it will not allow much interpénétration of surfaces and will not cause any numerical problem. By substituting 8f= 0.16 mm and Gc= 0.38 kJ/m in Eqn 1), a Tuit value of 4.75 MPa was obtained. This is initial assumption only, not the final value. ratio does not affect the overall delamination process much [10,11], With this advantage, crack

initiation displacement 8o was selected as 0.01583 to choose acceptable Keff value. By substituting initial T„it and on values in Eqn (3), we could get a reasonable Keff value of 300 MPa/mm and an Eeff of 300 MPa was obtained by Eqn (4). Effect of Kcff will be analyzed later for confirmation. All simulations were done with explicit mode of ABAQUS. So, mass density should be given in such a way that, it will not affect solution time increment of the problem and it should not impose much mass comparing bulk material. So, 10 % of the mass density value of PET was taken as mass density for cohesive elements. First, simulations based on surface based cohesive model were performed. The peel simulation input values & results for 0.2 mm long bulk mesh is shown m table III. Iterative simulations were started with 8f = 0.16 mm. The resulting peel force was higher than experimental force. Next the simulation was conducted with

S

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8f = 0.08 mm. The value was very high. So, it was decided to go with 8f value more than 0.16 mm. At 8f = 0.28 mm, the simulation results were matching closer with the experiment results.

No

1-A 1-B 1-C 1-D I-E 1-F 1-G 1-H

" Constar

Cohesive model

Surface based

Element based

ts: Bulk me

Table III.

Critical fracture energy

Gc (kJ / m2)

0.3800 0.3800 0.3800 0.3800 0.3800 0.1395 0.3800 0,1395

sh element lengt

Peel simulation input values & results for 0.2

Failure separation distance

or (mm)

0.08 0,16 0.20 0.24 0.28 0.28 0.28 0.28

ratio 0.4 0.8 1.0 1.2 1.4 1.4 1.4 1.4

Cohesive strength

T„,,(MPa)

9.5000 4,7500 3.8000 3.1667 2.7143 0.9964 2.7143 0.9964

mm bulk mesh.

Peel force from

simulation ( N )

13.50 11.10 10.82 10.62 10.49 3.61 10.93 3.93

FEA accuracy (%) 132.9 109.3 106.5 104.5 103.3 35.5 107.6 38.7

i, U =0.2 mm, Initial elastic stiffness, K eff =300 MPa/mm (Only for element based model), Cohesive mesh element length, lc= 0.04 mm (Only for element based model), lb/ lc ratio=5 ( Only for element based model )

Fig 6. Damage criterion for surface and element based models, (a) Peeling by surface based model. (b)View A-A: CSDMG for surface based model, (c) Peeling by element based model, (d) View B-B: SDEG for element based model.

Fig 7. 0.2 mm long bulk mesh peel test simulation by classical method. Fig 8. Reducing Tuli according to mesh in FEM. (a)Peel graph by surface model, (b) Peel graph by element model.

To cross check the result of surface based model, same simulation was done with element model with a value of Sf= 0.28 mm. Same stream No. I data from table II were used. Peeling progress by surface model is shown in Fig 6(a). In Fig 6(b), CSDMG value of 1 or above shows failed surface after delamination. Peeling by element model is shown in Fig 6(c), and in Fig 6(d), SDEG value of 1 or above shows failed cohesive elements after delamination. By seeing graph m Fig 7(a) it is clear that surface model predicts delamination well. Cohesive parameters related to 8f= 0.28 mm were fixed as final values for this mesh size. Similarly, graph of this same simulation No. 1-G based on element model is plotted in Fig 7(b). From results, it is understood that delamination prediction by both surface and element modeling are closer with minimum variation of 4 % approximately. To verify much plastic bending and deformation in peel arm, same simulations were done wjth cohesive parameters related to of value = 0.28 mm with ICPeel program Gc i.e., 0.1395 kJ/m by both surface and element model. This simulation values and results are given in table III (1-F, 1-H). By seeing the results, it is clear that simulation values are much lower than experimental ones. Hence, we can assume almost elastic deformation only had occurred in peel arm and if any plastic deformation occurred, it is negligible Peel test simulation using coarse meshes As said in introduction section, to use cohesive modeling in big size problems, the application of cohesive modeling should be verified in coarse mesh models. In element based cohesive model, A.Turon et al. [7] simulated double cantilever beam delamination using coarse meshes with both Gc and TU|t values known. He has shown that, by keeping Gc value as constant and reducing

294

interfacial strength from the original Tuit to reduced 7Z"* according to coarse mesh size by increasing the length of delamination zone (refer Fig Ï) will not affect the overall damage evolution significantly. Only the stress concentration at the crack tip will be lower. With these inputs, same 90 degree peel test simulations were conducted using coarse meshes. As in Fig 8, T„it value was reduced to "reduced cohesive strength" T%? by increasing 8f value. Bulk meshes of 1 mm and 2 mm long were simulated using surface based model. The initial configuration for 1 mm and 2 mm long mesh size were based on stream No. 2 and 3 from table II respectively. The Sf / lb of 1.4 corresponding to best simulation match with experiment in case of 0.2 mm bulk mesh length was maintained constant for both the cases. Now, with known Gc and of data, T'lf value can be calculated from Eqn (1). The simulation values and results are listed in table IV. Peeling progress, graphs are plotted in Fig 9(a) & 9(b) for 1 mm mesh and in Fig 9(c) & 9(d) for 2 mm bulk mesh. From graphs, it is clear that coarse mesh models also predict delamination with reasonable accuracy andthose results are closer to fine mesh models.

Table IV. Peel simulation input values & results for coarse bulk meshes.

No.

2-A 2-B 3-A

Cohesive Model

Surface based Surface based Element based

Failure separation distance ôf(mm)

1.40 2.80 1.40

Bulk material mesh length

lb (mm)

1 2 1

K ratio

1.4 1.4 1.4

Reduced Cohesive strength

/ ^ ( M P a )

0.5429 0.2714 0.5429

Peel force from

simulation ( N ) 10.39 10.11 10.59

FEA accuracy (%)

102.3 99.5 104.2

Constants: Ge= 0.3800 kJ/m , K eir =300 MPa/mm (Only for element based cohesive model), Cohesive mesh element length, lc= 0.2 mm (Only for element model), lb/lc ratio=5 ( Only for element model ), Average peel force from experiment = 10.16 N

Fig 9. Coarse simulation by classical energy method, (a) 1 mm mesh peeling by surface model, (b) Peel graph for 1 mm mesh peeling by surface model, (c) 2 mm mesh peeling by surface model, (d) Peel graph for 2 mm mesh peeling by surface model, (e) 1 mm mesh peeling by element model, (f) Peel graph for 1 mm mesh peeling by element model.

As done before, element based simulation also was done with 1 mm mesh for cross verification. Initial configuration was modeled based on stream No.2 from table II. The input values and simulation results have been given in table IV (No. 3-A). The peeling progress and peel force graph is shown in Fig 9(e) & 9(f). The average plateaus of peel force from both models are closer with less variation (3 ~ 5 % approximately). As discussed earlier, in element model, effect of Kcff value was verified by replacing initial value of 300 MPa/mm with 500 MPa/mm and 800 MPa/mm. For this study, 1 mm long bulk mesh case was taken. Simulation No. 3-A from table IV was simulated with kcffvalue 500 Mpa/mm and 800 MPa/mm. The peel forces predicted for ketr = 500 Mpa/mm is 10.68 N and for 800 MPa/mm is 10.71 N. These are almost closer to peel force predicted with Ke£f=300 MPa/mm since Ke„- was taken in reasonable range. Finally, comparison of approximate computational time was analyzed for both models. For this study, 0.2 mm and 1 mm length bulk mesh size models were run. In 0.2 mm long bulk mesh size, simulation no. 1-E from table III was taken for surface based cohesive model and simulation no. 1-G was taken for element based model. In 1 mm long bulk mesh size, simulation no. 2-A and 3-A were taken for surface and element based model respectively from table IV. All models were run twice and it was observed that the time taken by surface model is approximately 0.74 - 0.77 times of element based model in our case. In real life with bigger size problems, this much time saving would be a big one.

295

Conclusions The similarity between element and surface based cohesive model was taken in this study and element based cohesive constitutive response was applied in surface model. It was shown that closer prediction of delamination can be done with only two input cohesive parameters i.e., Gc, Tu]t using surface based cohesive model of ABAQUS. Surface based cohesive model predicts the delamination well closer to experiment. To verify possibility of using cohesive modeling in coarse mesh models, the same simulations were done with coarse meshes and it was shown that cohesive modeling can also be done with coarse meshes with modified cohesive parameters according to the required mesh size. To cross check the surface based cohesive model, the results obtained were compared with simulations based on element based cohesive behavior for sufficient cases. The simulations reveal that the results obtained by both methods were closer to each other with less variation ( 3 ~ 4 % ). Also, it was shown that surface model takes approximately 0.74 ~ 0.77 times of computational time compared to element based cohesive model in the present study.

Acknowledgment The sponsorship by Hyundai Motor India Ltd towards master degree program of author (Manivannan Ramamurthi) is acknowledged. The work was also supported by Priority Research Centers Program through NSF of Korea funded by MOEST(2011-0018392). Also the experimental support of Mr. Lee Jong Shin, Mr. Ko Sang Jin, Mr. Jeong Ji Yong of LG Electronics is acknowledged.

References 1. ABAQUS, ABAQUS Version 6.9-1 Documentation, ABAQUS, Inc, 2009. 2. X-P Xu, A. Needleman, "Void nucleation by inclusion debonding in a crystal matrix," Modeling and Simulation in Materials Science and Engineering, 1 (1993), 111-132. 3. G. Alfàno, "On the influence of the shape of the interface law on the application of cohesive-zone models," Composites science and technology, 66 (2006), 723-730. 4. K. Song, CG. Davila, C.A. Rose, "Guidelines and parameter selection for the simulation of progressive delamination," ( Paper presented at the ABAQUS Users' Conference, Newport, RI, 2008). 5. M.J. Van den Bosch, P.J.G. Schreurs, M.G.D. Geers, "On the prediction of delamination during deep-drawing of polymer coated metal sheet," Journal of Materials Processing Technology, 209 ( 2009 ), 297-302. 6. G. Alfano, M.A.Crisfield, "Finite element interface models for the delamination analysis of laminated composites: mechanical and computational issues," IntlJournalfor Numerical Methods in Engineering, 50 (2001), 1701-1736. 7. A. Turon, CG. Davila, P.P. Camanho, and J. Coasta, "An engineering solution for mesh size effects in the simulation of delamination using cohesive zone models," Engineering Fracture Mechanics, 74 ( 2007 ), 1665-1682. 8. J.G. Williams, H. Hadavinia, A.J. Kinloch, "Cohesive zone models in the characterization of toughness," ( Paper presented at the 11th International Conference on Fracture ICF11, Turin, Italy, March, 2005 ). 9. V. Gupta, A.S. Argon, D.M. Parks, J.A. Cornie, "Measurement of interface strength by laser spoliation technique," Journal of the Mechanics and Plastics ofSolids, 40 ( 1992), 141-180. 10. T. Diehl, "On using a penalty-based cohesive-zone finite element approach, Part I: Elastic solution benchmarks," International Journal of Adhesion & Adhesives, 28 ( 2008 ), 237-255. 11. T. Diehl, "On using a penalty-based cohesive-zone finite element approach, Part It: Inelastic peeling of an epoxy-bonded aluminum strip," International Journal of Adhesion & Adhesives, 28 ( 2008 ), 256-265. 12. I. Georgiou, H. Hadavinia, A. Ivankovic, A.J. Kinloch, V. Trospa, J.G. Williams, "Cohesive zone models and the plastically deforming peel test", The Journal of Adhesion, 79 ( 2003 ), 239-265. 13. D.R. Moore, J.G. Williams, "A protocol for determination of the adhesive fracture toughness of flexible laminates by peel testing; fixed arm and t-peel methods," (An ESIS Protocol, Revised June 2007, Nov 2010, http://www3.imperial.ac.uk/meadhesion/ testprotocols/ ) 14. A.J. Kinloch, C.C. Lau, and J.G. Williams, ICPeel 2006, "Gc Analysis for peel testing of Adhesives," ( Excel computer program, Imperial college, London, http:// www.imperial.ac.uk/me/ ) 15.H.F. Zhao, Wei, Yueguang, "Determination of interface properties between micron-thick metal film and ceramic substrate using peel test," International Journal of Fracture, 144 ( 2007 ), 103-112. 16. H.F Zhao, M. Chen, Y. Jin, "Determination of interfacial properties between metal film and ceramic substrate with an adhesive layer," Materials and Design, 30 (2009), 154-159. 17. O. Van der Sluis, P.H.M. Timmermans, E.J.L. Van der Zanden, J.P.M. Hoefhagels, "Analysis of the three-dimensional delamination behavior of stretchable electronics applications," Key Engineering Materials, 417-418 (2010), 9-12. 18. Laura De Lorenzis , Giorgio Zavarise, "Modeling of mixed-mode debonding in the peel test applied to superficial reinforcements," International Journal of Solids and Structures, 45 ( 2008 ), 5419-5436. 19. M.J. Van den Bosch, P.J.G. Schreurs, M.G.D. Geers, "Identification and characterization of delamination in polymer coated metal sheet," Journal of the Mechanics and Physics of Solids, 56 ( 2008 ), 3259-3276. 20. ASTM E8 /E8M - 09, "Standard test methods for tension testing of metallic materials," ASTM International, West Conshohocken, PA, United states, 2009. 21. ASTM D-882:2009, "Standard test method for tensile properties of thin plastic sheeting," ASTM International, West Conshohocken, PA, United states, 2009. 22. ISO 8510-1, "Adhesives- Peel test for a flexible-bonded-to-rigid test specimen assembly-Part 1: 90 degree peel".

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TIMIS2012 141 s t Annual Meeting & Exhibition

Nanocomposites

The proceedings contained in this section have not been edited or reviewed by the conference program organizers. The opinions and statements expressed in the proceedings are those of the authors only and arc not necessarily those of the editors or TMS staff. No confirmations or endorsements are intended or implied.

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

COMPRESSIVE STRENGTH OF EPOXY- GRAPHITE NANOPLATELETS COMPOSITES

H. A. Colorado1'2*, A. Wong', J. M. Yang1

'Materials Science and Engineering, University of California, Los Angeles. 2Universidad de Antioquia, Mechanical Engineering Department. Medellin-Colombia.

ABSTRACT

Epoxy matrix composites reinforced with Graphite Nanoplatelets (GNPs) at different concentrations have been fabricated in this research. The Epoxy matrix was made with Epon 828 cured with epikure 3055. GNPs were used in two conditions, as received and after being oxidized at high temperature. A planetary Thinky Mixing was used first to mix the Epoxy with the curing agent and then to mix the resulting liquid with the GNPs. The microstructure was identified by using optical and scanning electron microscopes and X-ray diffraction (XRD). Compressive strength was conducted over cylindrical samples of different length. X-ray micro tomography was used to characterize the samples after tested in compression in order to understand the failure modes.

Keywords: Polymer matrix composites. Graphite nanoplatelets. Mechanical properties.

'Corresponding author: Tel./Fax.: +1-310-206-8157/4830 Email: hcoloradolopera@uda.edu Address: MAE 38-137B, Engineering IV, 420 Westwood plaza, Los Angeles, CA 90095, USA

INTRODUCTION

Epoxy resin systems are increasingly being used as matrices in several industries including composite materials for structural [1-3], electronic [4,5], thermal [6,7] and aerospace applications. Other applications include casting resins [8,9], adhesives [10,11], and coatings [13,14] and tribological applications [15]. Epoxy polymers when cured are amorphous and highly-crosslinked polymers. This microstructure results in many useful properties for structural engineering applications, such as a high modulus and mechanical strength, low creep, and good performance at relatively elevated temperatures [16]. Unfortunately, due to the structure of the epoxy polymers, they are relatively brittle materials with a poor resistance to crack initiation and growth. The mechanical properties of epoxy matrices can be modified for several aspects, such as changing the molecular architecture and structure (i.e. by increasing the crosslink density to generate high stiffness and strength). Modifications can also be made by adding reinforcements such as a composite material with enhanced interfaces. It has been well established for many years that the incorporation of some second phase particles, (i.e. a thermoplastic micro polymer particles or silica nanoparticles [16]) can increase their toughness.

The aims of the present work are to investigate the strengthening of an epoxy polymer reinforced with graphite nanoplatelets (GNPs) and the damage after the compression test as well. A previous study [17] has shown promissory tensile results for epoxy resin with very low loadings

299

of GNPs.The sections below will present the compression strength and characterization with SEM and Micro-tomography for epoxy-graphite nanoplatelets composites. Samples were fabricated with GNPs as received and after an oxidation treatment. Tests were conducted for each GNPs type at 0.1, 1.0 and 5.0 wt% of particle loading. A reference sample without loading was also analyzed.

EXPERIMENTAL Samples Manufacturing In this research, the GNP was used as received and in an oxidized state. The GNP was oxidized by taking a glass beaker filled with the as received GNP and heated at 400°C for four hours; after being heated for four hours, the sample was slow cooled until it reached room temperature. The epoxy base was fabricated by mixing Epon 828 (resin) with Epikure 3055 (curing agent) in approximately a 1.7 ratio of resin to curing agent. Also, GNP was added to the mixture in different concentrations. In all cases, the resin and curing agent were mixed first; the mixing process normally took about 1 minute of manual stirring followed by 3 minutes of mixing in a Planetary Centrifugal Mixer (Thinky Mixer* AR-250, TM). Then, for samples with GNP, an additional 3 minutes of mixing in the Planetary Centrifugal Mixer was performed. After mixing, the mixture was poured into glass molds where it would cure at room temperature for 16-18 hours; it should be noted that mold release was applied at a minimum 30 minutes before having the mixture poured in. Finally, the glass molds were heat treated at 95°C for 2 hours. Compression Tests Samples for this test were fabricated using glass molds of 12.7mm diameter and 100mm long. A saw was used to cut cylinders of 15-16 mm in length. The samples were grinded to diameter length (approximately 11.5 mm when released from the glass molds) using a metallic mold until flat, parallel and smooth surfaces were obtained; silicon carbide papers of grit ANSI 200 and 400 were used. Compression tests were then conducted in an Instron machine. A set of 3 samples were tested for each composition at a crosshead speed of 1 mm/min. X-Rav Micro-tomography The equipment used was a Desktop Micro CT-System (Skyscan 1072), run at a high resolution mode. The X-Ray source was powered at 100 kV with a current of 100 uA and a power of 10W. No filter was used. For each sample, projections between 0.1° and 180° were taken in 0.25° rotation steps. Small camera pixel was used. The reconstruction was obtained from more than 300 individual x-ray radioscopes of the cross-sections using the Skyscan-NRecon Software. Other characterization Compression tests were conducted in an Instron 4411 machine. Samples were carefully polished to have two flat and parallel surfaces. Their dimensions were 11.5mm diameter and 11.5 mm length. The crosshead speed was 2.5 mm/min. Three samples were tests per composition. To see the microstructure, sample sections were progressively ground using silicon carbide papers of 500, 1000 and 2400 and 4000 grit. Once polished, these samples were first mounted on an aluminum stub and then sputtered in a Hummer 6.2 system at conditions of 15mA AC for 30 seconds to obtain a thin film of Au of around lnm. The SEM used was a JEOL JSM 6700R in a high vacuum mode.

ANALYSIS AND RESULTS

Figure la and lb display some of the samples before (top) and after (bottom) the compression tests were conducted for each composition. Figure 1 a shows samples fabricated with GNPs as

300

received. Figure lb shows samples fabricated with GNPs after thermal oxidation treatment at 400°C for four hours. Table 1 shows the samples dimensions before and after the compression tests. Figure lc shows SEM images of the GNPs as received. Figure Id shows SEM images of the GNPs after the thermal oxidation treatment. In general, after the thermal treatments, graphite layers on the GNPs seem to be moved (sliding one to each other) and less attached as well. Figure le and f show the composite with 5.0 wt% GNPs as received and after thermal treatment respectively. Figure le shows a cross section view image of a well impregnated GNP disc seen edge-on.

Figure 1 Manufactured samples before tested (top) after tested (bottom) for a) polymer coposites wuth GNPs as received, b) polymer composites with treated GNPs; SEM cross section view images for c) GNPs as received, d) GNPs after thermal oxidation, e) Polymer composite with 5.0wt% of untreated GNPs, f) Polymer composite with

5.0wt% of thermal treated GNPs.

Table I. Sample dimensions and tomography parameters.

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Figure 2a shows the mean compressive strength of all samples fabricated with their correspondent standard deviation. All samples' maximum compressive strength were over 120MPa. During the tests, no buckling was observed and sample was deformed symmetrically, which could be because the diameter to length ratio was kept at 1.0, the samples possessed flat and parallel surfaces and the high strength resin.

It is clear that the best results have been obtained for samples with 0.1 and 1.0 wt% of GNPs. When more GNPs were added, compressive strength decreased with respect to the sample with no GNPs (reference). The highest values were obtained for samples with 0.1 wt% oxidized GNPs (Ox GNP).

There was also a slight increase in the standard deviation for the samples with treated GNPs. This can be associated to the observation of less adhesion between graphite sheets for the GNPs after the thermal treatment. Under high compressive loads, the graphite sheets, which are stack on top of one another, easily slide past each other. Thus, when the GNPs content is higher, this effect is increased and less compressive strength is expected as shown in the results. Although it was not quantified in this research, porosity seems to have increased when GNPs content increased, which decreased the compressive strength. Figure 2b shows typical samples obtained for the composites and included in the data of Figure 2a. The stress-strain curve reflects the same tendency for all compositions respectively.

Figure 3 displays micro tomography cross section view images for polymer matrix-as received GNPs composites at the top, middle, and bottom portions of the sample's length. A side view x-ray image shows the sample. In all cases, the highest amount of damage was observed for samples at the middle of the sample.

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Figure 2 a) Compressive strength for the Epoxy-GNPs composite, b) typical stress-strain curves obtained.

On the contrary, the bottom portion of the sample possessed the least amount of damage. When compared with the reference sample (Figure 3a) at the middle, the sample's structural damage is reduced as the GNP content progressively increases.

Figure 4 shows micro tomography cross section view images for polymer matrix-thermal oxidized GNPs composites at the top, middle and bottom portions of the sample's length. A side view x-ray image shows the sample. Similar to the as-received GNPs composites, the highest amount of damage was observed for samples at the middle; the bottom portion received the least amount of damage. When compared with the reference sample (Figure 3a) at the middle, the sample's structural damage is reduced as the GNP content progressively increases. There is not an obvious difference in the fracture modes between samples with GNPs as-received and GNPs thermally treated.

SUMMARY

We have presented the compressive strength of epoxy-graphite nanoplatelets composites. All samples were over 120MPa at their maximum compressive strength. This is an inexpensive manufacturing method used for mixing without vacuum processing that showed satisfactory results when these composites were used for structural applications and high volumes. In one case, however, high temperature treatment was applied.

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Figure 3 Micro tomography cross section view images for polymer matrix-as received GNPs composites at the top, half and bottom of the length for a) reference sample, b) 0. lwt% GNPs, c) 1.0wt% GNPs and 5.0wt%GNPs.

The best compression strength was obtained for samples with 0.1 and 1.0 wt% of GNPs. When more GNPs were added, compressive strength decreased with respect to the sample with no GNPs (reference). In addition, there was a slight increase in the standard deviation for the samples with thermal treated GNPs. This can be associated to the observation of less adhesion between graphite layers for the GNPs after the thermal treatment. When the GNPs content is higher, this effect is increased and a lower compressive strength was expected as shown in the results. Although it was not quantified in this research, porosity increased when the GNPs content increased, which decreased the compressive strength.

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Finally, micro tomography shows than for the as-received GNPs composites, the highest amount of damage was observed for samples at the middle and the lowest damage at the bottom, for all samples. Also, it was observed in both cases that as GNPs content increases, structural damage is reduced. Additionally, not much difference was found in the fracture modes between samples with GNPs as-received and GNPs thermally treated.

Figure 4 Micro tomography cross section view images for polymer matrix-thermal oxidized GNPs composites at the top, half and bottom of the length for a) reference sample, b) 0.1wt% GNPs, c) 1.0wt% GNPs and 5.0wt%GNPs.

A C K N O W L E D G E M E N T S The authors wish to Colciencias from Colombia for the grant to Henry A. Colorado.

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3. B. Neffgen, "Epoxy resins in the building industry-25 years of experience," The International Journal of Cement Composites and Lightweight Concrete, 1 (4).

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5. Q.M. Jia, J.B. Li, L.F. Wang, J.W. Zhu, and M. Zheng, "Electrically conductive epoxy resin composites containing polyaniline with different morphologies," Materials Science and Engineering A, 448 (2007), 356-360.

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7. B. Perret, B. Schartel, K. Stoß, M. Ciesielski, J. Diederichs, M. Döring, J. Krämer, and V. Altstadt, "Novel DOPO-based flame retardants in high-performance carbon fibre epoxy composites for aviation," European Polymer Journal, 47 (2011), 1081-1089.

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10. Biqin Chen, Dongliang Jiang, Jingxian Zhang, Manjiang Dong, and Qingling Lin, "Gel-casting of ß-TCP using epoxy resin as a gelling agent," Journal of the European Ceramic Society, 28 (2008), 2889-2894.

11. A.H. Rezaifard, K.A. Hodd, DA. Tod and J.M. Barton, "Toughening epoxy resins with poly(methyl methacrylate)-gr tednatural rubber and its use in adhesive formulations," Int. J. Adhesion and Adhesives, 14 (2) (1994), 153-159.

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Micromechanical analysis of influences of agglomerated carbon nanotube interphase on effective material properties of a three-phase piezoelectric

nanocomposi t e

Tian Tang, Mark F Horstemeyer, and Paul Wang

Center for Advanced Vehicular System, Mississippi State University, Box 5405, Mississippi State, MS 39762, USA

Keywords: variational asymptotic method; three-phase piezoelectric nanocomposite; effective dielectric constants; agglomerated carbon nanotube interphase

Abstract

The focus of the present study is to investigate influ-ences of the agglomerated nanotube interphase on the effective material properties of a three phase piezoelectric nanocomposite using a recently devel-oped micromechanics framework, namely, variational asymptotic method for unit cell homogenization (VA-MUCH). The three phase nanocomposite is com-posed of PZT-7A fibers and epoxy matrix enhanced with single-wall carbon nanotubes (SWNTs). The effects of parameters of agglomerated nanotube in-terphase (caused by PZT fibers), such as properties and volume fraction, on effective dielectric constants of nanocomposite were analyzed.

Introduction

Piezoelectric materials generate voltage in response to an applied force, usually a uniaxial compressive force. Similarly, the application of voltage to a piezo-electric material will result in a change in its dimen-sions. Due to the drawbacks of pure piezoelectric materials such as low fracture toughness and high weight, piezoelectric composites were developed by dispersing piezoelectric materials in ductile materials such as polymers and metal alloys. Polymer matrix piezoelectric composites provide a greater piezoelec-tric response than conventional piezoelectric poly-mers as well as low density and high flexibility. How-ever, the polymer matrix piezoelectric composites are difficult to pole due to the large mismatch in the di-electric constants between the polymer matrix and piezoelectric inclusions. The low dielectric constants of the polymer hinder the applied electric field from adequately reaching the piezoelectric inclusions so that the composite response is small. On the other hand, the large electric field over piezoelectric inclu-sions caused by the externally applied force will be trapped by the polymer again due to the low dielec-

tric constants. One way to overcome this issue is to disperse the single-wall carbon nanotubes (SWNTs), which significantly increase the dielectric constants, in the polymer matrix to form three-phase nanocom-posites. However, as pointed out by Maxwell et al. [1], the piezoelectric inclusions will restrict the dis-persion of the SWNTs, resulting in a high density agglomeration of SWNTs that form a circular inter-phase region around the piezoelectric inclusions.

The focus of this paper is to investigate the in-fluences of the interphase on the effective dielectric constants of the composites using the recently devel-oped micromechanics model, Variational Ssymptotic Method for Unit Cell Homogenization (VAMUCH) [2,3, 4, 5, 6, 7]. This advanced model was constructed by invoking two essential assumptions: 1) exact solu-tions of the field variables have volume averages over the unit cell; 2) effective material properties are in-dependent of macroscopic geometry, boundary and loading conditions of the structure.

Piezoelectricity and Piezoelectric Composites

The elastic and dielectric responses are coupled in piezoelectric materials where the mechanical vari-ables of stress, and strain, as well as the variables of electric field and electric displacement, are related to each other. The coupling between mechanical and electric fields is described by piezoelectric coefficients. Using the conventional indicial notation, the linear coupled constitutive equations are expressed as:

(Jij = CijklCkl ~ CijkEk (1)

Ti = eiki€ki + kijEj (2)

where ay,-, e^, Ei and Ti are the stress tensor, strain tensor, electric field vector, and the electric displace-ment vector, respectively. Cijkt denotes the fourth-order elasticity tensor at the constant electric field,

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kij is the second-order dielectric tensor at the con-stant strain field, and djk is the third-order piezo-electric coupling tensor. To avoid the difficulty asso-ciated with heterogeneity, we can use the microme-chanics approach to homogenize the material and ob-tain an effective constitutive model, such that

Introduce the following matrix notation

7ij = c*jkieki -e*jkEk

Ti = e*jkckl + kfjEj

(3)

(4)

where "over-bar" indicates variables which are used in the macroscopic analysis of homogenized materi-als, and superscripts "*" denote the effective prop-erties whose calculations are determined by the mi-cromechanics model one employs.

The Variational Asymptotic Method for Unit Cell Homogenization (VAMUCH)

The derivation of VAMUCH starts from a variational statement of the heterogenous continuum. Taking advantage of the smallness of the microstructure, the homogenization problem can be formulated as a con-strained minimization problem posed over a single unit cell (UC) by carrying out an asymptotic analy-sis of the variational statement. The final theory of VAMUCH for homogenizing piezoelectric elastic ma-terials can be obtained by minimizing the functional

n„ Wn :TDedQ (5)

under the following periodic boundary conditions:

xV = X? and C " = Cr} for j = 1,2,3. (6)

In Eq. (5), e is a 9 x 1 matrix containing both the 3D strain field dj and the 3D electric field E{ and expressed as

€= [«11,2612, €22,2ei3,2e23, £33, EUE2, E3\T (7)

while D is a 9 x 9 matrix including the elastic, piezo-electric, and dielectric properties and is expressed as

-[-<-<] (8)

This constrained minimization problem can be solved analytically for very simple cases such as bi-nary composites. For general cases we need to turn to computational techniques for numerical solutions. Since VAMUCH theory is variational, the finite ele-ment method is a natural choice as a method to solve this problem.

a

¥ ôya 0 d

0 0 0 0

0

0 a

8a"

0 a

a»s

0 0 0

0

0 0

0 d ¥ ¥ 8» 0 0

0

0 o 0 0

0

0 a

"¥ ¥

= r « (9)

where Th is an operator matrix. If we discretize \ using the finite elements as

x(*<;w) = ■%<)*(*<) (10)

where 5 represents the shape functions and X a col-umn matrix of the nodal values of both the mechan-ical and electric fluctuation functions. Substituting Eqs. (9), and (10) into Eq. (5), we obtain a discretized version of the functional as

n n = ~(XTEX + 2XTDh€ë + ëTD«ë)

where E = j (ThS)TD(ThS)dü

Jil

■D« = f 1 Ja

Dh, 'Ddil

Ddü

(11)

(12)

(13)

(14)

Minimizing Hn in Eq. (11), we obtain the following linear system

EX = -Dheê (15)

It is clear from Eq. (15) that the fluctuation func-tion X is linearly proportional to c, which means the solution can be written symbolically as

X = XQC (16)

Substituting Eq. (16) into Eq. (11), we can calculate the electric enthalpy of the UC as

nn = 5£T (xjDhe + D.C) i = iTm Q

(17)

It can be seen that D in Eq. (17) is the effective piezo-electric material properties which can be expressed using a 9 x 9 matrix as

D = (18)

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and € is a column matrix containing both the global strains and global electric fields.

If the local field within the UC, such as local dis-placements, electric potential, stresses, and electric displacements, are of interest, they can be recovered in terms of their macroscopic behavior, including the global displacements Vi, the global electric potential ijj, the global strain and electric field e, and the fluc-tuation function x- First, the fluctuation functions should be uniquely determined. Considering the fact that we fixed an arbitrary node and made nodes on the positive boundary surfaces slave to the cor-responding negative boundary surfaces, we need to construct a new array XQ from XQ by assigning the values for slave nodes according to the corresponding active nodes and assign zero to the fixed node.

After having determined the fluctuation functions uniquely, we can recover the local displacements and electric potential as

121 I «3 |

'& dvi

ÖV2 OV1 OV% 3x7 3x7 5x7 flwa gya dvs dxt ox? 9x3 dip dif a y

{«} + SX

(19)

e = ê + ThSX (20)

a = Dt (21)

with a as a column matrix containing both 3D stresses and electric displacements such that

<7 — L(TH><T121<r22,ö'i3,<T23,0,33, ~Ti ,—T2, ~T3]

(22) The present theory has been implemented into a

companion code VAMUCH, a general-purpose mi-cromechanical analysis code. Although VAMUCH has all the versatility of the finite element method, it is by no means the traditional displacement-base finite element analysis. The code VAMUCH has the following distinctive features:

• No external load is necessary to perform the sim-ulation and the complete set of material proper-ties can be predicted within one analysis.

Figure 1: Unit nanocomposites.

cell of three-phase piezoelectric

Here S is different from 5 due to the recovery of slave nodes and the constrained node. The local strain field and electric field can be recovered as

Finally, the local stress and electric displacement field can be recovered straightforwardly using the 3D con-stitutive relations for the constituent material as

• The fluctuation functions and local displace-ments can be determined uniquely;

• The effective material properties and recovered local fields are calculated directly with the same accuracy of the uctuation functions. No post-processing type calculations such as averaging stresses and averaging strains are needed

• The dimensionality of the problem is determined by that of the periodicity of the UC. A complete 9 x 9 effective material matrix can be obtained even for a ID unit cell.

Resul ts and discussions

In this section, VAMUCH was employed to inves-tigate the influences of agglomerated carbon nan-otube interphase on the effective dielectric constants of three-phase PZT-7A fiber reinforced epoxy ma-trix enhanced by SWNTs. The UC of the three-phase nanocomposites is shown in Fig. 1. The PZT-7A fibers are of circular shape and in square array. Around the fibers is the interphase region composed of epoxy and agglomerated SWNTs. The matrix of nanocomposites is pure epoxy. The material prop-erties of PZT-7A and epoxy are taken from the lit-erature given by Berger et al. [8] as shown in Ta-ble. 1. The units of the elastic constants, piezoelectric constants, and dielectric constants are given in GPa, Cm - 2 , and nCV - 1m~1 , respectively. Since our goal was to study the influences of the interphase on the

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Table 1: Material properties of PZT-7A and Epoxy. constant interphase.

c„ Cl2 <?23 C22 C44

css Cm en e2i e3i esi fcn fc22 fc.33

PZT-7A 131.39 82.712 83.237 154.837

35.8 25.696 25.696

9.52183 -2.12058 -2.12058 9.34959

2.079 4.065 4.065

Epoxy 8.0 4.4 4.4 8.0 1.8 1.8 1.8

----

0.0372 0.0372 0.0372 Figure 2: Effective axial dielectric constants of the

three-phase nanocomposites as a function of volume fraction of the interphase.

effective dielectric constants of nanocomposites, the materials properties of interphase are set to the same as those of epoxy except that the values of dielectric constants vary with problems of interest.

Influence of volume fraction of interphase

At first, VAMUCH was used to investigate the influ-ences of the volume fraction of the interphase with a dielectric constant set as 372 nCV~ ' m - 1 . The effec-tive dielectric constants of the nanocomposites were calculated by changing the volume fraction of the interphase from 10% to 59% while the the volume fraction of PZT-7A fibers was kept constant at 20%. The effective axial and transverse dielectric constants were plotted in Figs. 2 and 3, respectively. In these two figures, the volume fraction 80% is a limit case at which the entire matrix has material properties of the interphase. It can be observed that the effective axial dielectric constants increase linearly as the vol-ume fraction of the interphase increases. The results illustrate that the effective transverse dielectric con-stants were negligible when the interphase was com-pletely enclosed by the low dielectric constant epoxy matrix in transverse directions and the volume frac-tion of interphase is less than 59%. They increased rapidly as the volume fraction of the interphase in-creased beyond 59%, after which the interphase was no longer completely enclosed by the low dielectric constant epoxy matrix. In other words, the effective dielectric constants increased significantly if the elec-tric field can be directly applied on the high dielectric

Figure 3: Effective transverse dielectric constants of the three-phase nanocomposites as a function of vol-uzne fraction of the interphase.

Influence of interphase properties

In order to investigate the influences of interphase properties on macroscopic material response, we eval-uated the effective material properties by increasing the dielectric constant of interphase from 0.372 to 372 n C V - 1 m - 1 while the volume fractions of fiber, interphase, and matrix were kept constant as 20%, 20%, and 60%, respectively. Fig. 4 and 5 show the variation of effective dielectric constants as the in-terphase's dielectric constant increased. We can ob-serve that the effective axial dielectric constants also

310

Figure 4: Variation of effective axial dielectric con-stants of the three-phase nanocomposites with the interphase's material properties.

Figure 5: Variation of effective transverse dielectric constants of three-phase nanocomposites with the in-terphase's material properties.

increased linearly with increase of interphase mate-rial properties. The effective transverse dielectric constants increased rapidly as the interphase's di-electric constant was less than 37.2. Then, they in-creased slightly as the interphase's dielectric constant increased further. Note that the effective axial di-electric constants were much larger than the effective transverse dielectric constants because the high di-electric constant interphase was completely enclosed by the low dielectric constant epoxy matrix in trans-verse directions.

Influence of volume fraction of PZT-7A fibers

In this section, the effective material properties were calculated by increasing the volume fraction of PZT-7A fibers up to 40% while the volume fraction of in-terphase was kept constant as 10%. The effective ax-ial dielectric constants and effective transverse dielec-tric constants were plotted in Figs. 6 and 7, respec-tively. Both effective properties increased slightly as the volume fraction of PZT-7A fibers increased be-cause the values of the dielectric constants of PZT-7A fibers are much smaller than those of the interphase. Since, in this case, the high dielectric constant inter-phase was completely enclosed by low dielectric con-stant epoxy matrix in transverse directions, the effec-tive transverse dielectric constants are much smaller than the effective axial dielectric constants.

Figure 6: Variation of effective axial dielectric con-stants of the three-phase nanocomposites with vol-ume fraction of PZT-7A fibers.

Conclusions

The advanced micromechanics model VAMUCH was used to evaluate the effective dielectric constants of the three-phase piezoelectric nanocomposites at dif-ferent parameters of the agglomerated SWNTs inter-phase. From this studies, it can induced that the re-sponse of nanocomposites is negligible if the electric field can not be directly applied to the high dielectric constant interphase. Therefore, it is necessary to in-crease the density of SWNTs in the epoxy matrix if the stronger response of nanocomposites is expected.

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Figure 7: Variation of effective transverse dielectric constants of the three-phase nanocomposites with volume fraction of PZT-7A fibers.

Acknowledgement

The authors acknowledged the Center for Advanced Vehicular Systems at Mississippi State University and the Department of Energy for supporting this research.

References

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[6] T. Tang and W. Yu. Variational asymp-totic micromechanics modeling of heterogeneous piezoelectric materials. Mechanics of Materials, 40(10):812-824, 2008.

[7] T. Tang and W. Yu. Asymptotical approach to initial yielding surface and elastoplasticity of metal matrix composites. Mechanics of Advanced Materials and Structures, in press, 2008.

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EFFECT OF NANO-PAPER COATING ON FLEXURAL PROPERTIES OF A FIRE-TREATED GLASS FIBER-REINFORCED POLYESTER COMPOSITE

Jamie Skovron, Jinfeng Zhuge, Ali P. Gordon, Jayanta Kapat, and Jihua Gou

Department of Mechanical, Materials, and Aerospace Engineering University of Central Florida, Orlando, FL USA

Keywords: Flexure testing, rupture, reinforced composite, thermo mechanical

Abstract

Planned re-usable aerospace vehicles require materials with high specific strength to withstand thermal shock associated with repeated re-entry into Earth's atmosphere. Composites, such as glass fiber-reinforced polyester (GRP), have rapidly become preferred for high value structural components requiring high specific strength and durability. Their ability to sustain high tensile loads, impact loads, and the like, has allowed them to be used as light-transmitting panels, fuselages, nose cones, and combustor nozzles. As a part of service conditions, heat flux strongly alters mechanical properties with exposure time. The effect of including a carbon nano-paper coating on the monotonie flexural properties of a GRP composite is analyzed. The nano-paper acts as a thermal barrier to protect the underlying material in the presence of above glass-transition temperatures. A series of three-point bend experiments was performed on specimen-sized samples of composites subjected to various levels of heat fluxes across numerous exposure times. Analysis of these experiments reveals trends in the deformation mechanisms of these materials near failure. Correlations of flexural modulus and critical load are used to develop models for strength.

1. Introduction

The use of composites has grown rapidly due to their attractive mechanical properties in comparison to existing conventional materials. Aircraft manufacturers have incorporated composites into more recent designs as a consequence. Frames of first-generation airplanes were constructed from wood, steel wire, and silk until it progressed to aluminum. This material has historically been the primary choice due to its low density, yet high tensile strength; however, more recently, the need for higher fuel economy has continued to alter aerospace materials selection [1]. The popularity of composites has grown tremendously in current commercial aircrafts. Components of new generation commercial aircrafts have grown from 12% to 50% by weight for composites; while aluminum has decreased from 50% to 20% by weight [2], Advanced research has demonstrated through testing, that composites have high specific strength at ambient temperatures (below 2\2°F) and thus have grown in reputation. This design philosophy is also applicable through development activities of re-usable launch vehicles. Such spacecrafts must be able to withstand the cold, near absolute zero, temperature occurring outside of the limits of Earth's atmosphere, but also endure the high temperature of approximately 1260°C (2300°^ during re-entry [3], The spacecraft components should be designed to endure these repetitive thermal exposures for several mission cycles while simultaneously being mechanically loaded, Re-entry vehicles that have traveled from Earth to Mars have encountered temperatures that exceed 1500°C (2732°F) [4]. It is essential that the selected fuselage materials

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be able to withstand these high temperatures without compromising structural integrity. With composites having such a high strength to density ratio, a range of 0.05 to 1 MJ/kg, they must be among the top competitors for material selection [1]. Design engineers must understand how composites will perform under certain loading and thermal conditions.

The focus of this study is to characterize how thermal barrier carbon nano-paper affects the flexural properties of GRP composites. Nano-paper coatings are applied to a GRP composite, which is then exposed to an applied heat flux and various exposure times. The heat flux is used to simulate the high re-entry temperatures. Following the return to room temperature (J2°F), flexural experiments are performed on the coated and uncoated samples. Results of the three-point bend test are outputted in the form of a force-displacement curve. The mechanical properties related to elastic, plastic, and rupture behavior of the samples are subsequently derived. Based on the experimental results and microscopy, analytical models are developed to predict the mechanical response of carbon nano-paper coated GRP composites under service conditions.

The candidate material for the current study is detailed in Sec. 2, and the experimental mechanics approach is described in Sec. 3. Section 4 of the current work focuses on the elastic properties of the GRP composite materials, while Sec. 5 covers observations on the rupture trends, respectively, that were developed. Conclusions and avenues for continued research are outlined in Sec. 6.

2. Glass-Fiber Reinforced Polyester (GRP) Composites

Glass-fiber reinforced polyester (GRP) composites, Fig. 1, are composed of multiple layers adhered by an orthophthalic resin. The polyester substrate material is reinforced with glass fibers whose specific strength is approximately 0.94 MJ/kg, this value gives the composite its high specific strength [5]. Use of GRP composites in more widespread applications is limited by their poor fire resistivity [6]; designers seek a thermal barrier to protect the effectiveness of the material. Carbon nano-paper is selected as the protective barrier that acts not only as a thermal barrier, but absorbs the heat that would otherwise penetrate the composite and compromise the structural integrity [7]. This study is focused on whether the carbon nano-paper can considerably protect the flexural properties of the material in the presence of high temperatures. Figure 2 details the dimensions of the uncoated and coated GRP composites.

Figure 1. GRP composite with flux induced side face up.

Figure 2: Dimensions of uncoated and coated GRP samples.

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3. Experimental Approach

The GRP composites were exposed to a predetermined flux level and exposure time as outlined in Table 1.

Table I: Heat treatment for the post fire 3 point bending test.

Applied Heat Flux (kW/m1) 25 35 50 75 100

Exposure Time (s) 0, 120,180,240,300 0,60, 100, 140, 180 0,40,80, 120, 150 0, 20, 50, 75, 100 0, 15,40,70, 100

Figure 3: Three-point bend test experimental setup

Mechanical experiments were then performed on samples of the candidate material to determine the mechanical effect of firing on the residual mechanical properties. Flexure experiments are preferred on composite samples since the imparted mechanical load bears resemblance to service conditions of the full scale components of these materials. Flexure tests were administered in accordance with the ASTM D790-10 standard [8]. A total of 117 samples underwent three-point bending, Figure 3, using an Instron 3369 with a load capacity of 50kN. The fire-treated side of the sample was placed face down during the three-point bending test. The support span, L, was 76mm and the crosshead motion rate, denoted by S was 2.4 mm/min. The experiment would cease if either the centerline deflection of the specimen reached 12mm or there was a dramatic drop in the load-deflection curve. The mechanical properties of these specimens such as modulus of elasticity in bending and rupture force can be calculated from the outputted raw data. By inputting variables into the model, the elastic modulus in bending (EB) was able to be calculated.

£ * ~ (1) * 4M3

Here, L is the support span (mm) between the two supporting rollers, m is the slope of the tangent to the initial straight-line portion of the load-deflection curve (kN/mm), b is the width of the beam tested (mm), and d is the thickness of the beam tested (mm).

4. Elastic Modulus

The elastic modulus in bending of the GRP composites was determined after the three-point bend test by utilizing Eq. 1. By inputting the parameters and data from the experimental procedure, the elastic modulus for each sample was calculated. The values for each flux-time sample set were averaged and used to formulate a model that predicts the elastic modulus based on a designated flux level and exposure time. Post-fire mechanical testing can hinder the true mechanical properties of the GRP composites during service. This deception is present due to the nature of composites as they tend to regain some of their strength after being cooled back down to room temperature [9]. For that reason, a temperature-time dependent model would be problematic to develop. Therefore, the following theorized model depicts the elastic modulus as a function of flux level and exposure time. Past researchers have developed mechanical property

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models based on separating the composite into two or three layers, but the drawback to these models is the classification of the different layers. [6,10]. The model presented in this study models the mechanical properties of the uncoated and coated GRP composites based on the flux level and exposure time. The model is to be utilized as a guide-line for the experimental data, any parameters outside the data set cannot be guaranteed for accuracy.

i=0 ;=0 (2)

In this equation, the flux level, 0, has units ofkW/m and the exposure time, texp, has units of seconds. The same model is used to estimate the elastic modulus for the uncoated and coated samples; however, the constants vary. The constants that must be inputted for either the uncoated or coated model are shown in Table II.

Table II: Constants for the elastic modulus model. Constant

Eoo E2i E01 En E10 E02 E,2

E20 E!3

Units

GPa s mN GPa/s 1/fjm

s//m \06Mg/msA

1 yms s'lm'g Am'x

Uncoated Value 23.28 5.802xl0'5

1.250x10'' - 8.545x10-' - 8.364xl0-2

0 0 0 0

Coated Value

20.06 4.210xl0J

8.219xl0-2

- 5.626xl03

-1.166x10-' 0 0 0 0

It was calculated that 93.75% of the uncoated data points fit within +/- 5.5GPa of the simulated values. The single data point that was not within this bound was the uncoated sample exposed to a flux level of 25 kW/m2 for 120 s. The value obtained experimentally was 2\.63GPa, but the simulated value was \4.90GPa.

Figure 4: Elastic modulus model for (a) uncoated and (b) coated GRP composites.

It was determined that all of the paper-coated data points fit within +/- 5.5GPa of the simulated values. In fact, there are only two data points whose difference between the simulated

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and experimental values is greater than 4GPa. These data points correspond to a heat flux of 25 kW/m2 and exposure times of 120 and 180 s. The R-squared value for the uncoated model is 87.7% and 80.8% for the coated one.

According to the data, regardless of <J) or constant texp, the elastic modulus decreases for either of these situations. The highest elastic moduli obtained during the experiment were the control samples for the uncoated and coated GRP composites. The control uncoated sample had an average value of 23.2SGPa, while the control paper-coated sample had a mean value of 19.52GPa. At the opposing end, the samples with the lowest elastic modulus values resulted from the greatest values of <)> and tcxp. The lowest elastic modulus obtained during the uncoated experiments belonged to the sample that had been exposed to a flux of 50 kW/m2 for a high exposure time of 120 s. The value was calculated to be as low as 0.895G/V», which is only 3.84% of the original elastic modulus obtained from the control sample. The coated sample that was exposed to a flux level of 100 kW/m2 for 100 s obtained the lowest calculated elastic modulus. This particular sample had a value of OJl&GPa, a value that is 3.99% of the elastic modulus from the paper-coated control sample.

5. Rupture

The effects of the addition of carbon nano-paper on the normalized maximum force and rupture behavior of the GRP samples are detailed based on the mechanical test data and the microscopy specimen pictures. Due to their nature, composites become pliable when exposed to a high temperature, thus the flexural properties measured after the heat flux process may not give accurate properties of how these composites act during fire [1 l].The characterization behavior of composites is a daunting task due to their anisotropic construction. This structural property alters their heat transfer process as when they burn, they release heat, particle-filled vapors, etc., which is followed by charring and finally delamination [11]. From the force-displacement curve, the maximum force applied during the three-point bend test was tabulated for each specimen sample. Three tests were administered for each combination of flux level and exposure time as shown in Table I.

A model was derived that predicts the normalized maximum force for the uncoated and coated samples using the input parameters of flux level and exposure time. The model is only valid for the data obtained during the three-point bend experiments; any values outside the data cannot be confirmed to be accurate. After inspecting the model, any normalized values that are less than 4% of the control specimen return a value that is less than 0. These samples are considered failed due to their low residual values. The constants, Fjj, are defined in Table 111.

The data for each flux level and exposure time combination were averaged and the trends were analyzed. In general it can be concluded that for each flux level set, the longer the exposure time, the lower the maximum force. This is due to the loss of the original material properties, and the material becoming 'charred' and rendered unusable in service conditions.

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Table III: Constants for normalized maximum force model

Constant Units Uncoated Coated Value Value 1.115 0.9983 5.779xl03 2.793x10-" 1.278x1 fr2 6.352xl0-3

4.575x10s 2.190xl08

-1.861x10"* -1.030xl0"6

-5.441X10"4 -2.793x10-" 0 0 0 0 0 0

Figure 5: Force-displacement for various data sets

The maximum forces for the uncoated and coated samples were normalized for each flux level. These normalized distributions are shown in Figure 6. One can see in Figure 6(a) that during a heat flux o f 25 kW/m2, the maximum force for the paper coated samples exposed to the flux for 180 s have almost three times the original maximum force compared to that of the uncoated samples, normalized values of 0.9598 and 0.3646 respectively. When the samples were exposed to a longer exposure time of 240 s, the normalized values became closer but the paper samples still maintained a higher percent of their original values, 0.222 compared to 0.029. The R-squared value for both the uncoated and coated model is 93 .3%.

Figure 6: Effects of exposure time on the normalized maximum force for (a) 25 kWm'ani (b) 35 kW m2.

The physical appearance of the GRP composites after being subjected to a level of heat flux and a three-point bend test is observed in this section. Figure 7(a) shows the composites exposed to a heat flux of 100 kW/m 2 with an exposure time of 70 s. The uncoated sample has continued to fray, in addition to the expansion of the fibers. This expansion is likely due to the gases escaping during the applied heat flux process [11]. The paper on the coated sample has become severely charred and has allowed a greater level of heat flux to penetrate the composite. The penetration of the coated sample is evident due to the increase thickness of the fibers. Fire barrier treatments, such as the carbon nano-paper used in this study, occupy the role of either reflecting the heat back towards the source or delaying the heat penetration towards the underlying composite [11].

Foo 1 F,„ m2,kW F„, / s F22 m2 (kN)2

F,2 -v (Mg) F„ m kN FOJ 1/s2

FM m"/MW2

F„ m2s/MN2

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Figure 7: GRP composite in cross-section after being exposed to a heat flux of 100 kW mr for 70 s then a three-point bend test for (al) uncoated and (a2) paper coated samples and a heat flux of 100 kWm2 for 100 s for (bl ) uncoated and (b2) paper coated samples. The left surface was exposed to the heat source.

Figure 7(b) is at the extreme end of the flux-time combination. The samples shown in Figures 7(bl) and (b2) have been exposed to a level of 100 kW/m2 for 100 s. Despite both samples having noticeably different physical appearances compared to their associated controls, the coated sample is still in one piece. Figure 7(bl), displays delamination of the composite.

Delamination is considered to be the sign of critical failure of a composite; it is characterized by fraying or ply separation due to interlaminar stresses [12]. This particular sample is to be considered beyond 'failed' due to it experiencing delamination and it being classified as a failed specimen before the three-point bend test. The coated sample was able to endure a three-point bend test, except it only retained a normalized maximum force value of 0.033, equating to 3.3% of the 15 s sample.

6. Conclusion

Glass fiber-reinforced polyester composites used for reusable launch vehicles can experience large declines in their mechanical properties after being exposed to a high heat source. The properties decrease as both the flux level and exposure time increases. The study confirms that the carbon nano-paper helps slow down the degradation process by acting as a thermal barrier between the heat source and the underlying composite. The data suggests that at high exposure times, the paper-coated samples were able to maintain a higher percent of the original elastic modulus compared to that of the uncoated samples. Models to simulate the elastic modulus and normalized maximum force were developed as a guide to understand the properties of GRP composites at various heat flux levels and exposure times.

Future work should be focused towards analyzing the coated and uncoated GRP composites under more extreme temperature values. The temperatures in space ranges from absolute zero to near 2500°F, thus a constitutive mechanical model must be developed for this wide temperature range. One must understand the characterization of these materials at extreme temperatures to ensure proper service.

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7. Acknowledgements

The materials presented here are based upon work supported by Office of Naval Research under Grant No. NOOO 14-09-1-0429 and Federal Aviation Administration Center of Excellence for Commercial Space Transportation (FAA COE-CST-AST) under grant number 10CCSTUCF002.

8. References

[ 1 ] Ashby, M. F., 2011. Materials Selection in Mechanical Design. 4th ed. Burlington, MA: Butterworth-Heinemann. pp. 12, Chap. 1.

[2] Hale, J. "Boeing 787 from the Ground Up." Aero Magazine. Vol. 4, pp. 17-24, 2006.

[3] Kulhman, E., "Investigation of high temperature antenna designs for space shuttle," Antennas and Propagation Society International Symposium. Vol.12, pp. 210- 213, 1974.

[4] Rawal, S., "Metal-matrix composites for space applications." JOM Journal of the Minerals, Metals and Materials Society. Vol. 53, Issue 4, pp. 14-17, 2001.

[5] Wambua, P., Ivens, J., Verpoest, I., "Natural fibres: can they replace glass in fibre reinforced plastics?", Composites Science and Technology. Vol. 63, Issue 9, pp. 1259-1264, 2003.

[6] Mouritz, A.P., Mathys, Z., "Post-fire mechanical properties of glass-reinforced polyester composites." Composites Science and Technology. Vol. 6, pp. 475-490, 2001.

[7] Zhuge, J., "Finite Element of Thermo-Mechanical Response of Fiber Reinforced Polymer Composites under Constant Heat Flux." MS. University of Central Florida, Orlando, FL, 2011

[8] ASTM Standard D790, "Standard Test Methods for Flexural Properties of Unreinforced and Reinforced Plastics and Electrical Insulating Materials," ASTM International, West Conshohocken, PA, 2010.

[9] Kandare, E., Kandola, B.K., Myler, P., Horrocks, A.R., Edwards, G., "Thermo-mechanical responses of fibre-reinforced epoxy composites exposed to high temperature environment: I." Journal of Composite Structures. Vol. 26, pp. 3093-3114, 2010.

[10] Bai, Y., Keller, T., Vallée, T., "Modeling of stiffness of FRP composites under elevated and high temperatures." Composites Science and Technology. Vol. 68, pp. 3099-3106,2008.

[11] Sorathia, U., Beck, C, Dapp, T., "Residual Strength of Composites during and after Fire Exposure." Journal of Fire Sciences. Vol. 11, Issue 3, pp. 255-70, 1993.

[12] Mallick, P. K., 2008. Fiber-reinforced Composites: Materials, Manufacturing, and Design, 3rd Ed, CRC Press, Boca Raton, FL. pp. 474, Chap. 6.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

FINITE ELEMENT MODELING OF THE NANOSCRATCHING OF POLYMER SURFACES

William M. Chirdon1 and Joshua T. Rozas1

'University of Louisiana at Lafayette; Department of Chemical Engineering 131 Rex St., Madison Hall; Lafayette, LA 70404, USA

Keywords: Polyethylene, Scratching, Nanoscratching, Finite Element Modeling

Abstract

Modeling of nanoscratching phenomena is important for a variety of material applications. In some applications, scratching is the desired effect of nanofabrication techniques. In other applications, scratching is undesired, as it causes coatings to fail, acts as a mechanism of wear, initiates cracks, and accelerates corrosion. This work uses finite element modeling techniques to investigate the effectiveness of various material models in their ability to model scratch experiments which have been previously published. The material models studied included viscoelasticity, deformation plasticity, elastic/plastic, strain-rate-dependent elastic/plastic, and other mechanical material models. Elastic/plastic models with and without strain-rate dependence were found to be the most robust material models. This work also highlights the importance of the scratching conditions, noting that the material properties that are desirable for resisting a scratching force are different from the desirable material properties for resisting a scratch displacement.

Introduction

A substantial body of experimental work has been accomplished by Misra et al in the investigation of scratching and nanoscratching of polymeric and metallic surfaces [1-6]. The objective of the presented work is to gain insight into the scratching phenomenon through finite element modeling. Specifically, the authors are modeling the nanoscratching of polymer and polymer composites that was accomplished by Yuan, Ramisetti, and Misra [7].

Selecting the most appropriate material model is an essential consideration for modeling scratching phenomena or other deformation behavior. In order to model any mechanical phenomenon, a mathematical model must be constructed to describe how a material will respond to a mechanical stress. However, modeling the response of a real material as a series of model equations can never be done perfectly, and selecting an appropriate material model is of paramount importance for the modeling of the nanoscratch. Various material models are fundamentally different in how they model a material's response to the stress, because they differ in how they include or neglect the dependence of response to factors such as strain hardening, strain rate, temperature, pressure, and other consideration related to the hysteresis of a material. The models also differ in how they define how much of the deformation is elastic (recoverable) and how much of the deformation is plastic (permanent). This poses a fundamental question for materials science, which is what is the best material model for modeling the polymeric response to a nanoscratch? In this study, the material properties are assumed to be homogenously defined throughout the sample. However, others have observed that the polymers may exhibit greater hardness in nano-scale deformations of their surface and may soften when indenting further into the sample [8],

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Related to this issue, is the importance of studying the effects of individual material properties or material model parameters on the nanoscratch response. For instance, if an engineer wishes to minimize scratching in an application, should the engineer seek a material with a high elastic modulus or a low elastic modulus? How is the scratch affected by the strain rate dependence of the material properties when scratched at different rates? Do the answers to these questions depend on whether a force or scratch depth is applied to the indenter?

Methods

The material model presented is composed of a combined elastic/plastic model. The elastic portion of the model considers the material to be a simple elastic solid defined by Young's elastic modulus and the Poisson's ratio. All of the deformation indictated by the elastic portion of the model is instantaneous and recoverable. The initial elastic model is defined by an elastic modulus of 200 MPa and a Poisson's ratio of 0.4. In series with the elastic model is the plastic model, which has been derived from uniaxial tensile stress/strain data. The plastic model consists of a plot of stress vs. plastic strain. Below the initial yield stress, there is no plastic strain, however, once the yield stress is exceeded, the amount of plastic strain quickly begins to exceed the amount of elastic strain. Figure 1(a) shows how the model fonctions as an elastic element and a plastic element (represented as a spring and dashpot) in series. Both elements share the same applied stress and their strains are additive. Note the dashpot element does not represent Newtonian fluid behavior, but instead a strain-dependent yield stress. Figure 1(b) shows stress/strain plot for the elastic and plastic models.

Figure 1. (a) Spring/dashpot diagram for elastic/plastic material. Note the dashpot represents a time-independent plastic strain in this specific model, (b) Stress/Strain diagram for the elastic and plastic material models.

Assembly of Finite Element Model A finite element model for the scratching of polymer surfaces has successfully been

assembled in ABAQUS and executed. Note that a biased mesh is used to maximize the accuracy near the scratch where a large number of nodal calculations are needed and a more coarse mesh is utilized further from the scratch in order to minimize the computational cost. In addition to the

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geometry, the boundary conditions, and the material model, the contact parameters between the indenter and the plate need to be specified. That is, without the definition of a contact to mutually exclude their volumes, the two objects will superimpose upon each other. The modeling of the contact of the sharp tip of the indenter as it moves across the surface has proven to be a substantial obstacle which has been overcome for the current model.

If the contact enforcement is too aggressive, it can cause excessive deformation of the surface elements which can cause the simulation to fail. Conversely, if the contact is not enforced effectively, the indenter will superimpose on the substrate instead of causing an indentation. It was found that the penalty method was the most robust method for efficiently enforcing the contact constraints with acceptable accuracy and computational cost.

To increase the efficiency of the calculation, the symmetry of the system is utilized. The scratching phenomenon should be symmetrical about the scratch, and so only half of the system needs to be evaluated and the other half can be assumed to be identical due to symmetry. Figure 2a) shows the actual model and nodes for calculation. Figure 2b) shows the model with its mirror image to help visualize the scratch. In this model, the indenter tip has a radius of curvature of 5 urn and an apex angle of 120°. The indenter moved at a scratch rate of 1 um/sec. Varying the scratch rates was not found to affect the scratch depth unless the material model is strain-rate dependent.

From the hundreds of permutations of material models, meshing techniques, and solution techniques that have been tested through this study, a selected subset of trials have been determined to be particularly informative and descriptive of the scratching phenomena. While various material models, including viscoelastic, Druker-Prager, Mohr-Coulomb, and deformation plasticity models have been tested, it has been concluded that a simple combination of the elastic and plastic models have been able to adequately model the scratching phenomenon. While the elastic/plastic combination model is the simplest of those tested, it has also been the most robust model which consistently converges to accurate results. Therefore, only the results of the elastic/plastic models are presented in this paper.

Figure 2. (a) Model for the scratch featuring half of the plate, (b) Visualization of the model results using a mirror image about the plane of the scratch.

Results

Effect of an Applied Force vs. an Applied Scratch Depth It was hypothesized that materials with higher elastic moduli would be more scratch

resistant than materials with lower elastic moduli. This hypothesis was validated by the current

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model in which all of the input parameters were held constant with the exception of Young's Modulus, which was varied. Note the plastic strain model, including the yield stress, was held constant in this series of experiments. An indentation force of 7 N was applied. It was observed that the instantaneous indentation depth increases rapidly as the modulus decreases and the material softens. In addition, there is a slight, yet significant, increase in the permanent indentation as the modulus decreases. This is due to the effect of the elastic response for a high-modulus material resulting in a higher elastic force that resists penetration in the lateral and vertical directions. However, an important finding has been that this hypothesis is only true if the scratch is created by an imposed force.

If the indentation force is not specified, but instead the vertical displacement (the depth of the indenter) is specified to be 200 um, the opposite trend is observed. That is, if the indenter moves across the material at a fixed depth, then materials with a higher modulus of elasticity will have a higher permanent scratch depth than a material with a lower modulus of elasticity. The permanent scratch depths observed for both boundary conditions are plotted as functions of the elastic modulus in Figure 3. This trend is found because the imposed strain for a higher modulus material results in a much higher stress. Since the yield stress is held constant, the higher stresses found in a higher modulus materials result in more plastic strain as the yield stress is further exceeded.

This leads to an important conclusion for applications where scratching is to be either performed or avoided. For instance, consider a biomédical implant in which scratching and wear is to be avoided. If this implant is pressed and sheared against a hard material with a given force, then a higher modulus material will be desired to minimize the amount of scratching and wear. However, if a computer-controlled scratching tool is used in the fabrication of a nano-device and a scratch is desired, and the tool is controlled to have a constant scratch depth, then the higher modulus material will result in a deeper permanent scratch.

Figure 3. Permanent scratch depth as a function of the elastic modulus under experimental conditions of applying 7 N offeree versus applying an indenter depth of 200 um.

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Similarly, the effect of the yield stress of the plastic model was investigated. In this experiment the stress of the plastic strain curve, including the yield stress, was scaled by factors of 0.5, 1, 2 and 4. As seen in Figure 4, modifying the plastic model stresses while holding the elastic modulus constant resulted in a sharp drop in the permanent scratch depth as the stresses of the plastic strain model were increased.

The results for single parameters mentioned thus far are of use and interest to materials engineers who model scratching. However, it is not realistic to assume that an engineer will have a variety of material choices with identical properties except for a range of various elastic moduli. More often, processes and compositions which increase the elastic modulus are likely to increase the yield stress. This is generally true for the strain hardening of metals and polymers, solid solution strengthening, or adding composite fillers to a polymer matrix. Therefore, the effect of simultaneously scaling the elastic modulus and the plastic model was investigated. Note that this will cause the depth of the scratch to decrease as the modulus and yield stress increase. The scratch depths as functions of the scalar factor of the plastic models with the same scalar factor applied to the elastic modulus is also shown in Figure 4.

Figure 4. Depth as a function of the scalar factor of the plastic strain with a proportional increase in the elastic modulus and with a constant elastic modulus.

Effect of Scratch Rate on Strain-Rate-Dependent Material Models Polymers are known for their heavily time-dependent properties. In order to incorporate

the time-dependent behavior of these materials, strain-rate dependent data was similarly decoupled into elastic and plastic models from experimental data published by Dasari and Misra [9]. The elastic model was defined by a various elastic moduli and a Poisson's ratio of 0.4. The model elastic strain was subtracted from the total strain in order to define the plastic strain as a function of stress.

When the material model does not have strain-rate dependence, the scratch depth was verified to be completely independent of scratch rate. However, once the strain-rate dependent

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models are implemented, the depth becomes scratch rate sensitive as shown in Figure 5. The effect on the horizontal displacement (in the direction of the indenter movement) is shown in Figure 6.

Figure 5. Permanent scratch depth as a function scratch rate for different moduli of elasticity using a strain-rate dependent plasticity model.

Figure 6. Permanent horizontal displacement as a function scratch rate for different moduli of elasticity using a strain-rate dependent plasticity model.

Investigation of Stick-Slip Phenomenon One discovery of interest was that when an applied indenter force (as opposed to vertical

displacement) was applied to a nano-indenter with a tangential friction coefficient of 0.3, the indenter would often bounce across the substrate as shown in Figure 7. This model utilizes a conical nano-indenter with a radius of curvature of 20 run and an apex angle of 65.3° with an

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applied force of 10 u.N. Originally, a few models have bounced very slightly over nodes, and this was presumed to be a numerical error due to the indenter crossing over a nodal integration point. However, in more refined models the bouncing would become more pronounced as shown in Figure 7.

The movement of the indenter creates waves in the material with a wavelength of approximately 4 micrometers. Several hypotheses were constructed to explain this bouncing behavior. The initial hypothesis that this was an error that existed about a node was refuted by the model shown in Figure 7, which includes deformations too large to be explained by a nodal anomaly. The next hypothesis was the indenter was bouncing on the substrate just as an idealized ball might bounce continually on an idealized elastic rubber material. This was originally supported by the observation that the indenter will bounce when the down-force is applied before the indenter moves horizontally. However, it was found that this type of bouncing is quickly damped in the model by plastic strain and the natural damping of numerical methods, whereas the observed bouncing does not become damped. Also, a purely elastic bouncing would have a defined period for bouncing; therefore, if the horizontal speed of the indenter was cut in half, die bounces would occur over half of the distance. It was found that changing the scratch rate did not impact the distance between bounces, which indicates that phenomenon does not arise from simple, periodic elastic bouncing, but rather represents the build up and release of shear stress at the surface due to frictional forces over a given distance. This phenomenon is very similar to the commonly observed ripples in the scratch path of micro-and nanoscratched polymers, which are proposed to arise from a 'stick-slip" mechanism as the frictional forces build up material and stresses beneath the indenter which are periodically released. This bouncing phenomena is presumed to be what was previously discussed in the founding literature as "Schallamach waves" which were originally observed by Schallamach [10,11]. This effect is seen in essentially all of the models in varying degrees, and poses difficulties in quantifying and comparing indenter depths across models.

Figure 2: Bouncing of indenter over substrate defined by an elastic/plastic material model.

Conclusions

Viable microscratching and nanoscratching models have been developed for testing material models. Implicit and explicit methods have been developed using various element types, and

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hexagonal C3D8R elements with distortion control on combined elastic/plastic material models using ABAQUS/Explicit has proven to be the most robust strategy for modeling nanoscratching phenomena. It was found that the underlying material properties desirable for scratch resistance depend on whether the scratch is imposed via a constant force or an applied depth. The use of strain-rate dependent data was found to be important for simulating experiments for different scratch rates. The bouncing observed in frictional studies is not a numerical error nor is it a simple elastic bounce from the material, but rather it is most likely the result of a stick-slip mechanism resulting in "Schallamach waves." All of these phenomena can be observed using the bulk properties of the polymer derived from uniaxial tensile testing data without utilizing spatially-dependent properties.

Acknowledgements

This project was sponsored by NSF Award #0852795. The authors wish to thank Ms. Natalie Girouard, Ms. Cristina Luna-Loya, Mr. Jeremy Leger, and Mr. Zachary Comeaux for their work on this project as undergraduate research assistants.

References

1. Dasari, J. Rohrmann, and R.D.K. Misra, "Micro- and nanoscale evaluation of scratch damage in polypropylenes," Macromolecular Materials and Engineering, 287 (2002) 889-903.

2. A. Dasari, S.J. Duncan, and R.D.K. Misra, "Atomic force microscopy of scratch damage in polypropylene," Materials Science and Technology, 18 (2002) 1227-1234.

3. A. Dasari, S.J. Duncan, and R.D.K. Misra, "Micro- and nano-scale deformation processes during scratch damage in high density polyethylene," Materials Science and Technology, 19 (2003) 239-243.

4. R. Hadal, A. Dasari, J. Rohrmann, and R.D.K. Misra, "Susceptibility to scratch surface damage of wollastonite- and talc-containing polypropylene micrometric composites, " Materials Science and Engineering A, 380 (2004) 326-339.

5. R.D.K. Mirsa, R. Hadal, and S.J. Duncan, "Surface damage behavior during scratch deformation of mineral reinforced polymer composites, " Ada Materialia, 52 (2004) 4363-4376.

6. N. L. Suryapudi, Q. Yuan, T.C. Pesacreta, and R.D.K. Misra, "The determining role of scratch indenter radius on surface deformation of high density polyethylene and calcium carbonate reinforced reinforced composite," Materials Science and Engineering A, 456 (2007)218-229.

7. Q. Yuan, N. Ramisetti, and R.D.K. Misra, "Nanoscale near-surface deformation in polymer nanocomposites," Ada Materialia, 56 (2008) 2089-2100.

8. S. Swaddiwudhipong, L.H. Poh, J. Hua, Z.S. Liu, K.K. Tho, "Modeling nano-indentation tests of glassy polymers using finite elements with strain gradient plasticity" Materials Science and Engineering A 404 (2005) 179-187.

9. A. Dasari, R.D.K. Misra, "On the strain rate sensitivity of high density polyethylene and polypropylenes," Materials Science and Engineering A 358 (2003) 356-371.

10. A. Schallamach, 'How does rubber slide,' Wear 17, (1971) 301-312. H . H . Tang, and D.C. Martin, 'Near surface deformation under scratches in polypropylene

blends,' Journal of Material Science, 38 (2003) 803-815.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

M A N U F A C T U R I N G A N D C H A R A C T E R I Z A T I O N O F AN A U X E T I C C O M P O S I T E

Fu-pen Chiang, Ph.D.

Department of Mechanical Engineering Stony Brook University

Stony Brook, NY 11794-2300, USA

Keywords: Auxetic Composite, Foam, Polymer, Speckle Photography

Abstract

Auxetic foam has a negative Poisson's ratio. As a result it exhibits some unusual mechanical properties that have advantages as a construction material especially as a ship building material. We manufactured an auxetic polymeric foam by 3D compression and thermal treatment of an ordinary foam. We then characterized its mechanical properties showing its advantages over the original material.

I. Introduction

An auxetic material has the unusual property of having a negative Poisson's ratio. When loaded in uniaxial tension it expands laterally rather than shrinks. This particular property gives rise to a variety of unusual mechanical behaviors when the material is subjected to load. Manmade auxetic material was first manufactured in 1987 by Lakes [1]. Subsequently, many investigators have ventured into the field [2-5]. An auxetic thin foam has several advantages over the conventional ones as a ship-building material. An auxetic thin plate deflects much less than a conventional plate for a given load [6]. It reduces acoustic noise due to its lower cut-off frequency [7]. It resists indentation and has a lower velocity impact damage [8, 9]. When bent it deforms synclastically rather than anticlastically thus rendering it ideally suitable for forming into convex-convex surfaces [10]. Furthermore, it resists shear failure due to the resulting large shear modulus. In this paper, we present the process of manufacturing an auxetic composite from an ordinary polyurethane foam. We characterize the mechanical properties of the resulting material.

There are different mechanisms that would result in auxetic behavior. One of which is the so called "re-entrant structure" whereby when loaded in tension the reentrant structure expands in transverse directions. An example is shown in Fig. 1 which is a microscopic picture of the auxetic polyurethane foams that we manufactured. Note the unit cell within the marked circle in the viewed area. The images depict a re-entrant structure. When loaded in tension vertically the re-entrant structure expands in both the vertical and horizontal directions progressively as the load is incrementally increased from picture # 1 to picture #4.

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Figure 1. Auxetic behavior in a foam due to "reentrant structure."

II. Manufacturing an Auxetic PVC Foam

We manufactured an auxetic foam by compressing an ordinary PVC foam (H45-Blue) in three dimensions. The PVC foam has the following properties: size of the foam cell: 50 um; tensile strength: -0.6 MPa; Young's modulus: 90Mpa; foam density: 44.3kg/m3. A 50x85x95 mm PVC foam block was first cut and then compressed in all 3 dimensions sequentially using an aluminum mold and a hydraulic press. After that, the aluminum mold was heated to about 135°C and then kept at that temperature for about 30 minutes. Finally, the aluminum mold was cooled down to room temperature and the resulting foam became auxetic. The global volume reduction of the fabricated foam was about 50% from the original but the local volume reduction varies. After compression and heat treatment cells become smaller and crushed inward. The schematic of the manufacturing process and micrographs of the original and auxetic foams are shown in Fig. 2.

Figure 2. Optical images of foam material (a) before processing, (b) after processing and the manufacturing scheme.

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III. Determination of Mechanical Properties of the Auxetic Foam

III-l Full Field Strain Measurement Using DSP (Digital Speckle Photography)

Foam is a relatively soft material for which a special full field displacement/strain measurement technique called DSP (Digital Speckle Photography) technique is employed [11]. The essence of the technique is as follows. A random speckle pattern of appropriate size is created on the surface (or interior) of the specimen. The speckle pattern before and after deformation of the specimen are digitally recorded. The digital images are divided into subimages of squared arrays of pixels (32x32, 16x16, etc., depending on the deformation of field) and then compared. The comparision is done via a 2-D Fourier transform algorithm called CASI [12]. The result is an array of displacement vectors each of which represents the aggregate displacement of all the speckles with the subimages. In some situations the surface texture of the specimen is itself a speckle pattern. And this is the case of the PVC foam we used. Thus, no treatment of the surface was needed.

111-2 Uniaxial Tension Test

Coupon specimens about 3.5mmx7mmx30mm were cut from an auxetic foam plate and tested in a uniaxial testing machine. A VH X-100 digital optical microscope was used to record the speckles (i.e., the texture variation of the specimen surface) as the specimen was being loaded quasi-statically at 2mm/minute. A typical stress-strain curve is shown in Fig. 3.

Figure 3. Typical stress-strain curves of auxetic foams

This results in a Poisson's ratio of about (-0.16). A full field displacement map of the specimen under load is depicted in Fig. 4. When the displacement vector distribution clearly indicated that the entire specimen expands as the tensile load is being applied. Additional tests have shown that

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the degree of auxeticity of the foam is a function of the volume contraction ratio during the manufacturing process. The larger the volume reduction, the larger the resulting auxeticity.

Figure 4. Example of displacement fields obtained by CASI. v field, u field and total displacement vector. Contour Constant is 31m

III-3 Indentation and Low Velocity Impact Tests

Because of the negative Poisson's ratio, an auxetic material tends to shrink laterally when subject to compression. Thus it tends to resist indentation and low velocity impact loads. For the indentation test, we used an Instron machine with a small spherical loading tip to load both the auxetic and ordinary foams. Fig. 5 shows two indentation marks formed on the materials. It is seen that the indentation mark of the ordinary foam at 35N is much larger than that of the auxetic foam at 60N.

Figure 5. Indentation area under the view of microscope Left: Conventional Foam, IA: d=12mm at 35N; Right: Auxetic foam, IA: d=7.5mm at 60N

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Figure 6. Penetration values of the metal bullets into the foam block as a function of Poisson's ratio

A low velocity impact test was carried out using a Daisy 105B air gun with 4.5mm steel bullets. Both the foam and the gun were fixed on a table top with the gun being about 4 feet away from the foam specimen. The penetration depth of both the auxetic foam and the ordinary foam were measured and plotted as shown in Fig.6. It is seen that with the same impact energy, the penetration depth of the auxetic foam is much larger than that of the ordinary foam.

Hl-4 Cyclic Loading Test

We also performed various cyclic loading tests of both the auxetic and ordinary foams. The load was applied quasi-statically at the speed of 2mm per minute. Fig. 7 shows some of the results. It is interesting to note that at the first two cycles the stress-strain loop traverses a counter clockwise direction. After that the stress-strain loop reverses its direction and traverses clockwise and this is what one would expect in a traditional material. We also performed a test on the original foam with identical loading conditions. The loop always rotates clockwise. This is easily explainable in that the energy stored in the loading stage (the area under the stress-strain curve) is always less than the energy recovered when the load is released due to the loss as a result of internal friction. In the case of the auxetic foam, the original foam cells are crushed into such a shape that they form re-entrant corners. And these corners retain a fair amount of elastic energy. At the beginning of the tensile loading cycle this elastic energy is recovered as the reentrant corners are being straightened. Thus when the load is reversed, this energy is recovered adding to the total energy recovered giving rise to the unusual phenomenon that it appears that one recovers more energy than one puts in. Subsequent loading depicts the conventional cyclic stress-strain loop indicating that the material is no longer auxetic.

We also continued the cyclic loading until the specimen breaks. We find that the stiffness of the auxetic foam decreases rather significantly as the loading cycle increases. Fig. 8 depicts the degradation of the material stiffness as a function of loading cycles for two auxetic foams (with Poisson's ratio being (-0.08) and (-0.14) respectively.

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Fig 7 Stress-strain loops of the auxetic foam (a) Loading force from 0-10N-0; (b) Loading force from 0-15N-0; (c) Loading force from 0-20N-0; (d) Loading force from 0-25N-0; (e) Loading

force from 0-30N-0; (f) Loading force from 0-35N-0

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Figure 8. Stiffness degradation of auxetic foams as a function of loading cycle

In Fig 8, Fo is maximum load F obtained in the first cycle and F is maximum load obtained in the subsequent cycles. The auxetic foam was broken at an average number of cycles 920 (results vary from 810 to 1180), compared to the conventional foam which was broken at an average number of cycles 360 (results vary from 230 to 430). The comparison shows that the auxetic foam has a higher resilience towards fatigue damage as well.

IV. Conclusion and Discussion

We have demonstrated that a conventional PVC foam can be converted into an auxetic foam with a negative Poisson's ratio by 3-D compression and heat treatment. The auxeticity starts at about 50% reduction in volume of the original foam. The larger the volumetric reduction (up to 70% in the current study), the higher the auxeticity. We show that the resulting auxetic foam has better resistance, as compared to the original foam, to indentation and low velocity impact damage. It seems to have better fatigue life as well.

V. Acknowledgement

We would like to thank Dr. Yapa D.S. Rajapakse of ONR for supporting this work through grant No.N000140410357.

VI. References

1. R.S. Lakes, "Foam structures with a negative Poisson's ratio, " Sei, 235 (1987), 1038-1040.

2. A. Alderson, and K.E. Evans, "Molecular origin of auxetic behavior in tetrahedral framework silicates," Phys Rev Lett, 89 (22) (2002), 225503-1.

3. K.L. Alderson, A. Fitzgerald, and K.E. Evans, "Strain dependent indentation resilience of auxetic microporous polyethylene," J of Mater Sei, 35 (16) (2000), 4039-4047.

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4. K.L. Alderson, A. Alderson, and K.E. Evans, "Interpretation of the strain-dependent Poisson's ratio in auxetic polyethylene," J of Strain Anal for Eng Des, 32 (3) (1997), 201-212.

5. R.H. Baughman, J.M Shacklette, et al., "Negative Poisson's ratios as a common feature of cubic metals," Nature, 392 (6674) (1998), 362-365.

6. K.E. Evans, "Auxetic polymers: a new range of materials," Endeavour New Series, 15 (4) (1991), 170-174.

7. C.P. Chen, and R.S. Lakes, "Dynamic wave dispersion and loss properties of conventional and negative Poisson's ratio polymeric cellular materials," Cell Polym, 8 (5) (1998), 343-359.

8. F. Scarpa, J.R. Yates, et al., "Dynamic crushing of auxetic open-cell polyurethane foam," Proc of the Inst ofMech Eng, Part C: J ofMech Eng Sei, 216(12) (2002), 1153-1156.

9. F. Scarpa, L.G. Ciffo, et al., "Dynamic properties of high structural integrity auxetic open cell foam," Smart Mater and Structures, 13 (1) (2004), 49-56.

10. J.P.M. Whitty, A. Alderson, P. Myler, and B. Kandola, "Towards the design of sandwich panel composites with enhanced mechanical and thermal properties by variation of the in-plane Poisson's ratios," Compos Part A: Appl Sei and Manuf, 34 (6) SPEC (2003), 525-53.

l l .F.P. Chiang, "Evolution of white light speckle method and its application to micro/nanotechnology and heart mechanics," Opt Eng, 42 (5) (2003), 1288-1292.

12. D.J. Chen, F.P. Chiang, Y.S. Tan, H.S. Don, "Digital speckle-displacement measurement using a complex spectrum method," Appl Opt, 32 (11) (1993), 1839-1849.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MICROTRUSS CELLULAR NANOCOMPOSITES

Khaled Abu Samk1, Guojie Huang2, Milan Skocic3, Hatem S. Zurob2, D. Embury2, Olivier Bouaziz4, Glenn D. Hibbard1

University of Toronto - Materials Science and Engineering; 184 College Street; Toronto, ON M5S 3E4, Canada

McMaster University - Materials Science and Engineering; 1280 Main Street West; Hamilton, ON L8S 4L8, Canada

^Grenoble Institute of Technology; 7 Rue Massena; Grenoble, 38000, France

4ArcelorMittal Research; Maiziere-les-Metz, 57283, France

Keywords: Microtruss, Carburizing, Martensite, Inelastic Buckling

Abstract

New types of cellular nanocomposites can be created by carburizing the external surface of low carbon steel microtruss materials. Microtruss architectures are designed to resist externally applied loads through axial deformation and as such can exhibit significantly enhanced strength and stiffness when compared to conventional metal foams. They can also exhibit multifunctional characteristics such as energy absorption and thermal management. This study used the high formability of simple low carbon sheet steels to create an initial microtruss cellular architecture by plastic deformation. Mechanical strength was imparted via a graded composite structure wherein an external skin of ultrafine internal length scale martensite was created by carburizing. The microtruss nanocomposites exhibited up to a five-fold increase in compressive strength when compared to the conventional low carbon steel microtruss reference. The failure mechanisms were investigated in order to determine an optimal composite structure.

Introduction

There is a growing need for lightweight materials that can reduce energy consumption in the transportation and aerospace sectors. Microtruss cellular materials are designed such that externally applied loads are resolved along the internal struts making up the cellular architecture, allowing them to exhibit high strength and stiffness at low density [1]. Microtruss materials can be made using simple sheet forming methods [2, 3]. Using plain carbon steels would be particularly attractive, because these materials are produced on a massive scale and have lower embodied energies (energy consumed during manufacturing) compared to other lightweight alloys of aluminum, magnesium, and titanium [4]. However, while low carbon steels exhibit excellent formability [5], their strength is usually low.

One way to increase the strength of a low carbon steel microtruss is to electrodeposit a coating of nanocrystalline Ni around the internal struts [6]. Although this sleeve represents additional mass, the strength increase associated with having nm-sized grains in the coating [7], combined with the optimal positioning of this coating away from the neutral bending axis of the struts, results in a specific strength increase for the nanocomposite [6], On the other hand, heat treatable steels have a rich potential to create ultra-high strength nanostructures through conventional deformation and phase transformations [8]. One approach to increasing the strength of low

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carbon steel microtrusses would be to carburize the surface after fabrication in order to induce the formation of ultra-fine length scale martensite laths at the outer surfaces - in effect creating a composite structure having an ultra-high strength outer shell and ductile inner core. In addition, this approach would not suffer from the weight penalty issues arising from reinforcement by electrodeposition of nanocrystalline Ni. The present study examines the feasibility of using carburizing as a strengthening mechanism in low carbon steel microtruss cores.

Experimental

Low carbon steel (AISI-SAE 1006) sheets with square perforations were purchased from McNichols Perforated Products (Atlanta, GA). See Fig. la for a schematic diagram of the perforated sheet with the dimensions labeled. The perforated sheets were annealed at 800°C for 15 minutes to relieve strain-hardening and internal stresses induced while rolling and perforating the sheets. The annealing process is followed by air cooling to room temperature then the sheets are formed into pyramidal microtrusses using a mechanical press that stretches and bends the internal struts to form the final architecture, after [6, 9]. A combination of three forming and annealing steps were used to create the final architecture shown in Fig. lb.

Fig. 1 Schematic diagram of (a) the starting low carbon steel perforated sheet (L = 6.36 mm, I, = 5.10 mm, w/ = 1.26 mm, t, = 0.64 mm), and (b) the final geometry of the low carbon steel microtruss after forming (H = 5.82 mm, 6 = 46°, / = 7.34 mm, w = 1.02 mm, t = 0.58 mm).

Four sets of microtrusses (four repeats each) were produced using the same fabrication technique, with the only difference being the final heat treatment (Table 1). Samples were tested in uniaxial compression at a cross-head displacement rate of 1 mm/min using confinement plates to prevent displacement of the microtruss nodes, thus inducing the same inelastic buckling failure mechanism that would occur in a fully bonded sandwich panel, after [10].

Table 1. Final heat treatments applied to four subsets of fabricated low carbon steel microtrusses. Subset A

B

C

D

Symbol WH

AN30-AC

AN30-WQ

C180-WQ

Final Heat Treatment Work hardened condition. No heat treatment after last forming cycle; imparted plastic deformation during final fabrication step is retained. Annealed for 30 minutes at 800°C to remove fabrication-induced work-hardening, followed by air cooling to room temperature. Annealed for 30 minutes at 830°C, followed by water-quenching to room temperature. Carburized for 180 minutes at 830°C in CO/CO2 gaseous mix at a ratio of 4.69 computed using ThermoCalc to achieve a target carbon content of 0.4 wt.%C, followed by water-quenching to room temperature.

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Results and Discussion

The martensitic phase transformation is a diffusionless process that occurs during the quenching of low- or medium- carbon steel from the austenite region. The carbon atoms in the interstitial sites of the face-centered-cubic (fee) austenite are trapped upon quenching, resulting in a military transformation of the austenitic fee phase into the martensitic body-centered-tetragonal (bet) phase, resulting in the formation of ultra-fine length scale martensite laths aligned parallel to one another within areas of the parent austenite grain [11], While individual martensite laths are too small to be seen by optical microscopy, the martensite regions can be clearly seen in the polished and etched cross-sections of sample types C and D. Fig. 2a shows the small fraction of martensite formed from the original carbon content in the AIS11006 alloy, while Fig. 2b shows the formation of a hardened case surrounding the strut after carburization and the formation of multiple martensite islands in the core due to the diffusion of carbon through the grain boundaries. This microstructural gradient can also be seen from microhardness profiles through the strut cross-section (Fig. 3), where the hardness of the outer case is up to five times higher than the hardness of the inner core.

Fig. 2 Strut cross-sections of (a) C: AN30-WQ sample, and (b) D: C180-WQ sample. The light region is ferrite while the dark region is martensite.

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Fig. 3 Vickers hardness profile for the carburized subset D: C180-WQ. The dotted horizontal line is the average reference hardness before carburizing.

Carburizing also results in a dramatic change to the uniaxial compression strength of the microtruss cores. Fig. 4 presents typical stress-strain curves for each of the four sample types outlined in Table 1. In each case, there was an initial elastic region followed by a load drop as the individual struts began to buckle. The fully annealed and air-cooled samples (subset B) had a peak strength of 2.83 ± 0.04 MPa, which represents the base strength of the low carbon steel microtrusses. In contrast, sample type A had a peak strength of 4.63 ± 0.12 MPa, indicating that the work-hardening introduced during the final fabrication step plays a major role in determining the strength of the as-formed samples. While -40% of this strength is lost after recrystallization and air-cooling from 800°C (i.e. going from sample condition A to B in Table 1), nearly the same magnitude of strength increase is recovered by the formation of martensite islands (Fig. 2a) after water-quenching from 830°C (the peak strength of sample type C was 4.59 ± 0.31 MPa). However, the largest change was seen after carburizing followed by water-quenching. In subset D, the microtrusses were heated to a temperature of 830°C in a gaseous mixture of CO/CO2 in order to increase the surface carbon content from a max of 0.08 wt.%C [12] to 0.4 wt.%C. This treatment formed a hardened case on the surface of the microtruss, increasing the strength to 14.45 ±0.11 MPa, i.e. more than a five-fold increase over the strength of the sample type B reference. These peak strength values are summarized in Fig. 5a and the relative significance of these strengthening mechanisms are illustrated in Fig. 5b.

Fig. 4 Stress-strain curves for the four microtruss subsets tested in uniaxial compression.

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Fig. 5 (a) Peak strengths of the microtruss subsets A through D. (b) The individual columns indicate the strength added by the different strengthening mechanisms observed in this study.

While carburizing resulted in a significant strength increase, it also changed the post-buckling behaviour of the composite microtrusses. This could be seen in the stress-strain curves of sample type D where small load drops could be detected after the initial peak strength was reached. Fig. 6 plots the local tangent modulus of the uniaxial compression stress-strain curves as a function of strain for the four different sample types. Sample types A, B, and C show a relatively smooth transition in tangent modulus from peak to valley (minimum) strength, while sample type D exhibits sharp transitions in tangent modulus after the peak strength has been reached. Sharp load drops were also seen in nanocrystalline Ni reinforced Al microtrusses and were correlated to delamination and crack propagation through the nanocrystalline sleeve [13].

Fig. 6 The slope of the tangent to the stress-strain curves plotted as a function of strain. Sharp deep valleys in the carburized microtrusses (subset D) represent cracking/fracture events.

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In order to examine the post-buckling failure mechanisms, a series of low magnification SEM images of samples loaded to beyond the peak strength were taken in order to examine the overall shape of the struts and to detect surface cracks formed during testing. Fig. 7a shows that samples from subset C (and likewise subsets A and B) exhibit the typical morphology expected of inelastic buckling, and no evidence of crack formation could be detected. In contrast, Fig. 7b shows that the samples from subset D, while also exhibiting inelastic buckling, had surface cracks in the regions experiencing local tensile stresses (e.g. the convex side of the buckling strut). These cracks reduce the load carrying capacity of the post-buckled struts and contribute to the steep drop in strength of the microtruss beyond the initial peak (Fig. 4).

Fig. 7 Inelastically buckled struts of (a) subset C: AN30-WQ showing no cracking after buckling, and (b) subset D: C180-WQ showing crack formation after buckling.

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As the sample is further loaded beyond the initial peak, cracks of the type seen in Fig. 7b eventually become complete fractures. SEM was used to take characterize the fracture surfaces of these failed composite struts. Fig. 8a shows a complete fracture of an individual strut in a microtruss from subset D. Fig. 8b shows that crack propagation is influenced by the microstructure shown in Fig. 2b. As the strut continues to deform past the peak, the tensile stresses on the outer surface of the buckled strut increase resulting in the formation of cracks through the hardened case. Once the cracks have reached the ferrite phase in the strut, they become blunted as the ferrite phase can easily deform. However, at higher strains, as the cracks reach the martensite islands in the core, they propagate more easily, connecting other martensite islands and leading to the formation of the smooth cleaved regions seen in Fig. 8b.

Fig. 8 (a) Lower magnification, and (b) higher magnification, SEM images of fracture surface of individual strut in a carburized microtruss from subset D: C180-WQ crushed past the peak.

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The results presented in this study indicate that carburizing can be used to significantly enhance the strength of the low carbon steel microtrusses. The 3-hour carburizing treatment used in this study resulted in a five-fold increase to the inelastic buckling strength of the low carbon steel microtrusses, without the accompanying weight penalty that would be associated with nanocrystalline electrodeposition. The fact that load drops and crack propagation were only detected beyond the initial peak, suggests that the strength of the composite microtruss (sample type D) is controlled by the same type of inelastic buckling failure mechanism as the three reference sample types (A, B and C). However, in order to minimize crack propagation and enhance the post-buckling load carrying capacity of the composite microtruss, it may be necessary to trade-off a certain amount of peak strength increase from carburizing, either by tempering the composite microtruss after quenching or else by applying an additional decarburizing step that would create a thin ductile ferrite shell around the hardened martensite case.

References

[I] VS. Deshpande, M.F. Ashby, and N.A. Fleck, "Foam topology bending versus stretching dominated architectures," Ada Materialia, 49 (6) (2001), 1035-1040. [2] H.N.G. Wadley, "Multifunctional periodic cellular metals," Philosophical Transactions of the Royal Society a-Mathematical Physical and Engineering Sciences, 364 (1838) (2006), 31-68. [3] D.J. Sypeck and H.N.G. Wadley, "Cellular metal truss core sandwich structures," Advanced Engineering Materials, 4(10) (2002), 759-764. [4] T.E. Norgate, S. Jahanshahi, and W.J. Rankin, "Assessing the environmental impact of metal production processes," Journal of Cleaner Production, 15 (8-9) (2007), 838-848. [5] ASM International, Metals Handbook, vol 1 - Properties and Selection: Irons, Steels, and High-Performance Alloys (Materials Park, OH: ASM International, 1990). [6] B.A. Bouwhuis et al., "Structural nanocrystalline Ni coatings on periodic cellular steel," Composites Science and Technology, 69 (3-4) (2009), 385-390. [7] N. Wang et al., "Room temperature creep behavior of nanocrystalline nickel produced by an electrodeposition technique," Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 237 (2) (1997), 150-158. [8] B. Chehab et al., "Bulk nanoscale materials in steel products," in I5th International Conference on the Strength of Materials (ICSMA 15), Dresden, Germany, 2010. [9] B.A. Bouwhuis and G.D. Hibbard, "Failure Mechanisms during Periodic Cellular Metal Fabrication by Perforation Stretching," Metallurgical and Materials Transactions a-Physical Metallurgy and Materials Science, 39A (12) (2008), 3027-3033. [10] B.A. Bouwhuis and G.D. Hibbard, "Compression testing of periodic cellular sandwich cores," Metallurgical and Materials Transactions B-Process Metallurgy and Materials Processing Science, 37 (ISSUE) (2006), 919-927. [II] G. Krauss, Steels: Heat Treatment and Processing Principles. Materials Park, Ohio: ASM International, 1990. [12] ASTM International, Standard A830/A830M-11: Standard Specification for Plates, Carbon Steel, Structural Quality, Furnished to Chemical Composition Requirements. [13] E. Bele, B.A. Bouwhuis, and G.D. Hibbard, "Failure mechanisms in metal/metal hybrid nanocrystalline microtruss materials," Ada Materialia, 57 (19) (2009), 5927-5935.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

SYNTHESIS OF TAILORED CORE-SHELL MAGNETIC MICROPARTICLES FOR INTRAVASCULAR EMBOLIZATION

Gabriella R. Ferreira, Alexandre P. Umpierre, Fabricio Machado*

Institute de Quimica, Universidade de Brasilia, Campus Universitäre Darcy Ribeiro, CP 04478, 70910-900, Brasilia, DF, Brasil - E-mail: fmachado@unb.br

Keywords: Intravascular embolization, Hyperthermia, Vinyl acetate, Magnetic nanoparticles

Abstract

Spherical core-shell microparticles of magnetic poly(vinyl acetate) / poly(vinyl alcohol) polymers intended for medical applications such as intravascular embolization procedures and hyperthermia treatment of injured areas were synthesized by in situ incorporation of surface modified magnetite nanoparticles through suspension polymerization process. The proposed experimental methodology showed to be very efficient for the synthesis of micro-sized magnetic particles with controlled morphology. Partially hydrolyzed magneto-polymeric nanocomposites exhibited good thermal stability that observed for the pure poly(vinyl acetate). It was also noted that the morphology of the precursor polymer is preserved when mild reaction conditions are applied during the hydrolysis reaction. Micro-sized polymeric particles showing good magnetic properties were obtained, which indicates that core-shell poly(vinyl acetate) / poly(vinyl alcohol) polymers can be successfully used as embolie agents.

Introduction

In the recent years, the development of new classes magneto-polymeric particles intended for applications in biotechnology and biomédical areas of research has become an increasingly important field. Magnetic particles has been widely used in several kind of application purposes, as for instance, nanoprobes for imaging in vivo, contrast agents in magnetic resonance imaging, immobilization of enzyme, magnetizable implants for targeted drug delivery, stimuli-responsive systems, electromagnetic wave absorbers, hyperthermia treatments, environmental analysis and magnetic separations, and catalytic action of magnetic nanocomposites [1-10].

Intravascular embolization is a technique widely used as surgical intervention procedure for the treatment of tumors, aneurysms and arteriovenous malformations. This technique can be distinguished by the injection of finely divided material, via catheter, into the bloodstream around the tumor region in order to mechanically obstruct the vessels responsible for supplying the blood to the injured area, interrupting the supply of oxygen and nutrients to the tumor region, which helps the reduction of tumor size, allowing for the recovery of tissue after short time interval [11-17],

Poly(vinyl alcohol)-based embolie agents present some desirable characteristics such as, high compressibility and good elasticity, excellent biocompatibility and high thermal stability. The combination of these main features normally ensures satisfactory sterilization of the polymer, an easy insertion and transportation of the compressed material through the catheter, proper blood vessel occlusion, minimization of morphological changes due to temperature effect, and good resistance to acids, alkalis and detergents [12, 18].

It is generally agreed that the morphology of embolie agents plays a fundamental role on the performance of the intravascular embolization procedure, since the undesirable inflammation

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of the embolized vessel walls and vascular recanalization, which minimizes the success of the blood vessels occlusion, are closely linked to irregular morphology of embolie agent particles. Problems associated to poor morphological features of embolie materials can be generally overcome when spherical polymer particles with narrow size particle distribution are produced.

In this scenario, new embolie agents can be developed with a structural concept based upon the core-shell type structure, characterized by the combination of a rigid domain (due to the thermoplastic polymer present in the nucleus) and controlled flexible phase (polymer with good compressibility in the spherical particle shell). The structure-performance relationship is very dependent on the material employed in manufacture of both the core and the nucleus of the polymer. Based on this particle design, poly(vinyl alcohol)-based embolie agents can be successfully produced presenting a controlled malleability restricted only to the shell, due to the thin layer of poly(vinyl alcohol), PVA that reduces an eventual recanalization of the blood vessels and inflammation in the walls of vascular tissues caused by irregularities in size and shape of embolie particles, while the core comprises the precursor polymer formed by poly(vinyl acetate), PVAc, which is an attractive material largely studied because of the excellent optical properties, good biocompatibility [19].

The development of polymeric materials exhibiting magnetic properties arises as a very attractive possibility to the improvement of the characteristics of traditional PVA / PVAc embolie agents. The dispersion of magnetic nanoparticles into the thermoplastic matrix allows the effective use of an external magnetic field to help the transportation of the magnetic embolie material at specific points in the human body, and additionally this new material can be successfully used for the hyperthermia treatment and prevention in the injured area subjected to embolization procedure [20].

This work focuses on the synthesis of magnetic particles of poly(vinyl acetate)-based polymeric materials produced through suspension polymerization process by in situ incorporation of surface modified superparamagnetic Fe3Û4 nanoparticles intended for in vivo biomédical applications of intravascular embolization.

Experimental

Materials Nitrogen was supplied by White Martins Ltda, Rio de Janeiro, Brazil, with 99.5% purity. Distilled water was used as the reaction medium. The reagents sodium hydroxide (NaOH) with purity of 99%, ferric chloride hexahydrate (FeCl3-6H20) with purity of 97% and ferrous sulfate heptahydrate (FeSC>4-7H20) with purity of 99% were provided by VETEC Quimica Fina Ltda (Rio de Janeiro, Brazil), hydrochloric acid (HC1, 36.5-38.0% w/w) was provided by 1SOFAR (Sào Paulo, Brazil) and oleic acid extra pure was provided by Merck (Rio de Janeiro, Brazil). Vinyl acetate (Vetec Quimica Fina Ltda - Rio de Janeiro, Brazil) was used as monomer. The suspending agent [poly(vinyl alcohol) (PVA), DENKA POVAL B-24] with a hydrolysis degree of 86-89% was kindly donated supplied by DENKA, Tokyo, Japan. The initiator of polymerization [benzoyl peroxide (BPO), LUPEROX 78] with a minimum purity of 99.4% was kindly donated by Arkema Quimica Ltda, Säo Paulo, Brazil. Sodium sulfate anhydrous (Na2S04) with purity of 99%, acetone [(CH3)2CO] with purity of 99.5%, methanol (CH3OH) with purity of 99.8% and ethanol (C2H5OH) with purity of 99,8% were provided by Vetec Quimica Fina Ltda (Rio de Janeiro, Brazil). Sodium dodecyl sulfate with purity of 99% was provided by Quimibrâs Indüstrias Quimicas SA (Rio de Janeiro, Brazil). All chemicals were used as received, without further purification.

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Synthesis of Magnetite Nanoparticles Magnetite nanoparticles of Fe3C>4 were synthesized based on chemical coprecipitation method of aqueous Fe2+ and Fe3+ salt solutions and NaOH solution, as described elsewhere [21]. In order to obtain black magnetic particles, it is necessary to guarantee the stereochemistry of the reaction by adopting a molar ratio of Fe2+ in relation to Fe3+ equal to 2.0. The coprecipitation reaction was carried out as follows: 6.1 g of FeClv6H20 and 3.1 g of FeS04-7H20 were dissolved in 5mL of hydrochloric acid (used to help the solubility of the salts in water), followed by dilution in 125mL of distilled water and heated at 60 °C with a nitrogen gas flushing. In another reaction vessel, 37.5 g of NaOH were dissolved in 625mL of distilled water, and heated to 60 °C with purge of nitrogen gas. The salts solutions were added to the basic solution and kept under vigorous mechanical stirring. The temperature of the reaction medium was maintained at 60 °C for 30 minutes under atmosphere of nitrogen to preserve the reaction environment free of oxygen and prevent oxidation of nanoparticles. The resulting black magnetic nanoparticles were separated by décantation employing a magnetic bar of neodymium and washed with distilled water until neutral pH. Finally, the FejC>4 nanoparticles were both washed several times with and stored in ethanol.

Surface Modification of Magnetite Nanoparticles The modification of nanoparticles surface with oleic acid is fundamental to ensure good dispersion inside the monomer microdroplets during polymerization reaction. The modification reaction was performed based on the following procedure [22, 23]: Initially, 5 g of magnetic nanoparticles were dispersed in 170 mL of distilled water with mechanical stirring, a continuous bubbling of nitrogen gas and temperature heating until 85 °C. After that, 5.5 mL of oleic acid were added dropwise at a constant rate of 0.5 mLmin"1 under an inert nitrogen atmosphere. After complete oleic acid addition, the reaction system was kept under mechanical stirring at 85 °C for 30 minutes. The Fe304 nanoparticles with modified surface were washed several times with distilled water to obtain pH near 7, and then washed three times with acetone. Finally, the resulting surface modified Fe304 were dispersed in vinyl acetate for polymerization.

Polymerization The polymerization reactions were performed in a 500-mL glass reactor (Quickfit, England) under a nitrogen atmosphere in order to keep the reaction environment inert. The reaction vessel was equipped with a hotplate 1KA C-MAG HS 7 (IK.A Works, Inc.), connected to an integrated temperature control and a PtlOOO temperature probe, was used to maintain the temperature of the reaction medium at the setpoint; a reflux condenser, coupled to a cold water feed stream; and an overhead Stirrer RW 20 digital (IKA® Works, Inc., Säo Paulo, Brazil) equipped with a helix-type impeller. Before starting reaction, the reactor was fed in with 200 mL of a solution containing 0.235 g/L of PVA (suspending agent) and the system was heated to 65 °C with constant agitation of 500 rpm. After the settled temperature was reached, a mixture containing 0.4 g of BPO, 8 g of surface-modified Fe304 nanoparticles dispersed in 80 g of vinyl acetate was added in the reactor under agitation. The reaction was maintained at 70 °C during the first two hours and the raised to 75 °C in the last two hours in order to maximize monomer conversion. At the end of the reaction, the polymer was washed with a solution of sodium dodecyl sulfate to eliminate a residual monomer.

Partial hydrolysis Reactions of Polymers Hydrolysis reactions of both pure poly(vinyl acetate) - PVAc and poly(vinyl acetate) magnetic nanocomposites - PVAc / Fe3Û4 were carried out in test tubes with aqueous solutions of NaOH/CH3OH and NaOH/CH3OH/Na2S04 in different concentrations, based on the following

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procedure: in each test tube, it was added 20 mL of a solution containing NaOH/CfbOH or NaOH/CH30H/Na2S04 and 1 g of PVAc or PVAc / Fe3C>4. The suspension was maintained under magnetic stirring at desired temperature (50 °C or 70 °C) and reaction time (ranged from 4 to 36 h). At the end of this reaction, the polymer was washed with distilled water until pH 7.

Material Characterization The materials were characterized by different experimental techniques. Particle morphology and particle size distributions analyses of Fe3C>4 nanoparticles were performed by transmission electron microscopy - TEM (Jeol, JEM-1011, Tokyo, Japan). Energy dispersive X-ray spectroscopy (EDX) measurements of Fe3C>4 nanoparticles and PVAc / Fe3Û4 nanocomposites were performed on an EDX-720 fluorescence spectrometer (Shimadzu Europa GmbH, Duisburg, Germany). The XRD patterns of the Fe3C>4 nanoparticles and PVAc / Fe3Û4 nanocomposites were obtained on a Bruker D8 FOCUS X-ray diffractometer (Bruker AXS, Inc., Wisconsin, USA), using CuKa radiation (X = 1.5418 Â, 40 kV and 30 mA) operating in the angular range 2 < 26 < 80 with steps of 0.05° at a rate of 0.25°-min"'. The surface morphology of polymeric particles was observed by scanning electron microscopy (SEM) on a FEI Quanta 200 Scanning Electron Microscope (FEI Company, Oregon, USA). Thermo gravimetric measurements were carried out in order to evaluate the thermal stability of materials on a Thermogravimetric Analyzer Shimadzu DTG-60 (Shimadzu Scientific Instruments, Maryland, USA) at heating rates of 10 °C-min"', under nitrogen atmosphere and flow rate of 50 mL min"1. The thermal transitions of the polymeric materials were analyzed by DSC measurements in a Shimadzu DSC-60 calorimeter (Shimadzu Scientific Instruments, Maryland, USA) at heating rates of 10 "Cmin"'.

Results and Discussions

Figure 1 shows the typical morphology of magnetic nanoparticles observed by transmission electron microscopy (TEM). As can be seen, the Fe304 nanoparticles present a narrow particle size distribution with average size (dp ) around 8.4±2.4 nm. This characteristic size indicates that Fe3Û4 nanoparticles synthesized behave like a superparamagnetic material, which is strongly recommended for in vivo biomédical applications such as intravascular embolization procedures intended to treat tumors, arteriovenous malformations (AVMs) and aneurysms, as the magnetic retention takes place only when an external magnetic field is applied. As a matter of fact, magnetic nature of the material depends significantly on the size of the magnetic nanoparticles.

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Figure 1. TEM of Fe304 Magnetic Nanoparticles (A) and Particle Size Distribution of Fe304

(B). Figure 2 shows the diffractogram of Fe304 nanoparticles. It can be observed the presence

of typical peaks of magnetite in 26 = 18.6, 30.2, 35.5, 43.5, 53.6, 57.3, 62.7 and 74.6°, corresponding to following reflections (111), (220), (311), (400), (422), (511), (440) e (533), indicating that the synthesized magnetite has a crystalline structure of the type of spinel [24-27]. The crystallite size (ÇNp) of the Fe304 nanoparticles determined based on the Scherrer's equation [N. B. in order to calculate ÇNp, the X-ray wavelength (X) was adopted to be equal to 0.15418, the constant k equal to 1.0 and the (311) reflection was used to determine the full width at half maximum (FWHM)] was estimated to be equal to 7.6 nm, which agrees very well with the average particle size obtained by TEM measurements.

Figure 2. X-ray Diffraction Patterns of Fe304 Magnetic Nanoparticles

The thermal stability of polymeric materials and the weight fraction of magnetic nanoparticles dispersed into the polymeric matrix of pure PVAc and core-shell PVAc / PVA were evaluated by thermogravimetric measurements, as shown in Figure 3. The total degradation of the polymeric materials is observed within temperatures ranged from 280 °C to 750 °C, which demonstrates that the polymers present good thermal stability. Typically, two important weight-losing take place: the first occurs at temperatures in the range from 280 °C e 370 °C (corresponding to 60 wt-% of weight loss) and the second significant weight loss occurs in the range from 400 °C to 750 °C, when the complete thermoplastic matrix disappears. Differences in the patterns of thermal decomposition (TD) are noticed for the final temperature of decomposition between PVAc/PVA equals to 750 °C and the precursor PVAc, whose TD corresponds to 600 °C, making evident the improvement of thermal stability of the final material that is appropriate for vascular embolization. According to Figure 3, weight fraction of Fe304 nanoparticles was determined to be equal to 10 wt-% for both PVAc / Fe304 and PVAc / PVA / Fe304 nanocomposites. Based on the weight loss results, it seems that the leaching of Fe304 nanoparticles does not occur during the hydrolysis of the precursor PVAc / Fe304.

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Figure 3. Thermal Stability of Polymers Materials Evaluated by Themogravimetric Analysis

Polymeric particles of PVA / PVAc with good morphological features are obtained. As illustrated in Figure 4A, the polymer particles exhibit some macrocavities, probably due to release of encapsulated water into the thermoplastic matrix of polymer during the polymerization of PVAc. The surface of the particles is also characterized by presenting roughness, probably due to the combined effect of methanol excess and collision between the particles as a result of the agitation of the reaction medium when the hydrolysis reaction of PVAc is performed. In the particular case of Figure 4B, particles presenting smooth surface is obtained due to the effect of Na2SÜ4 on the solubility of the polymer particles. In this case the Na2SO,t provides double benefit: firstly, it minimizes the solubility of the polymer chains of PVAc in methanol, and the secondly, it reduces the solubility of the polymer chains of PVA in aqueous media.

Figure 4. Morphology of Polymer Particles of PVA/PVAc.

Figure 5 illustrates the morphology of magnetic particles of PVAc and PVA / PVAc. As shown in Figure 5A, spherical particles with good morphological characteristics can obtained. On the other hand, particles of PVAc / Fe304 exhibit a characteristic porosity, which may appear as a consequence of the modified Fe3Û4 dispersed in the polymer matrix of PVAc. It is possible that this feature of the porous particles of PVAc / Fe3Û4 allows a more effective action of the methanol solution on the solubility of the PVAc polymer chains, leading to the formation of

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particles of PVA / PVAc / Fe3C>4 with irregular morphology, as shown in Figure 5B. The beneficial effect of the salt is exemplified once again, as the incidence of particles with irregular morphology is minimized when the hydrolysis of the PVAc / Fe304 is carried out in the presence of Na2S04, as illustrated in Figure 5C.

Figure 5. Morphology of Magnetic Polymer Particles of PVA/PVAc.

Conclusions

Magnetic polymeric particles with controlled morphology were successfully obtained. The partial hydrolysis of PVAc leads to an increase in the thermal stability of final product. It was also observed that the morphology of precursor polymer was preserved when mild reactions conditions regarding the system temperature and alkoxide concentration were adopted during the hydrolysis. In addition, it was exemplified that the use of Na2SC>4 is strongly desirable for morphologic control purposes, acting decisively over the decrease of the solubility of PVAc by the methanol and PVA by the water.

Acknowledgement

The authors thank Coordenaçao de Aperfeiçoamento de Pessoal de Nivel Superior (CAPES), Conselho Nacional de Pesquisa e Desenvolvimento (CNPq) and Fundaçao de Empreeendimento Cientificos e Tecnolögicos (FINATEC) for the finnantial support and the scholarships. Particularly, F. M. thanks Arkema for providing benzoyl peroxide and DENKA for providing Poly(vinyl alcohol), and Laboratorio de Materiais Combustiveis (LMC), Laboratorio de Quimica medicinal e Tecnolôgica (LaQuiMeT), Laboratorio de Quimica Analitica e Ambiental (LQAA) and Laboratorio de Microscopia Eletrônica e Virologia for the research suport.

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Jolly EE. J. Vase Intervent Radiol 2008;19(l):42-46. 14. Kerber CW, Bank WO, and Horton JA. Amer J Roentgenol 1978;130(6). 15. Hurst BS, Stackhouse DJ, Matthews ML, and Marshburn PB. Fertil Steril 2000;74(5):855-

869. 16. Peixoto LS, Cordeiro FB, Melo PA, Nele M, and Pinto JC. Macromol Symp 2011;299-

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DRAMATIC EXPANSION OF LUMINESCENCE REGION IN GaP/POLYMER NANOCOMPOSITES

Sergei L. Pyshkin1'2, John Ballato2

'institute of Applied Physics, Academy of Sciences, Academy St. 5, Kishinev, MD2028, Moldova 2Center for Optical Materials Science & Engineering Technologies and the School of Materials Science and Engineering, Clemson University, Technology Dr. 91, Anderson, SC, 29625, USA

Keywords: Luminescence, GaP, Polymer, Nanocomposites

Abstract

GaP nanoparticles were prepared using mild aqueous or colloidal syntheses at decreased temperature followed by ultrasonication and stored as dried powder or suspension in water-ethanol mixture or toluene. Selected mixtures of GaP nanoparticles, based on dimensions, exhibit at room temperature bright broad band luminescence from UV until yellow-red region with controlled band-width and spectral position of maximum. Poly (2-vinylpyridine) (P2VP), biphenyl vinyl ether (BPVE) and tetrahydrofuran (THF) polymers were used to prepare GaP based nanocomposites. The resulting nanocomposites showed pronounced quantum confinement effects and other discussed in this work important for application interesting phenomena leading to dramatic I eV expansion of GaP luminescence to the UV spectral region. These GaP nanoparticles and GaP/polymer nanocomposites provide significant enhancement of blue-shifted luminescence from which novel light emissive device structures may be fashioned.

Introduction

While bulk and thin film GaP have been successfully commercialized for many years, its application in device nanocomposite structures for accumulation, conversion and transport of light energy has only received attention recently. Increasing interest to some GaP/polymer nanocomposites is explained by complementary behavior of components in some composites where each constituent is a candidate for use in light emitters, waveguides, converters, accumulators and other planar, fiber or discrete micro-optic elements. In the present paper we discuss preparation and properties of the GaP nanoparticles and light emissive films on the base of some GaP/polymer nanocomposites. This work has been fulfilled in framework of the STCU (www.stcu.int ) 4610 Project "Advanced Light Emissive Device Structures" and continue our efforts, discussed at the 2006, 2010 and 2011 Nanotech Conferences, 2007-2011 TMS Annual Meetings, Conferences and Symposia, and published in the relevant proceedings and papers [1-14] with the focus being to advance the quality and light emissive properties of GaP nanocomposites and using our new results in preparation of closed to ideal bulk GaP single crystals [2-8, 13, 14], different methods of GaP nanoparticles syntheses and the most optically and mechanically compatible polymers.

Experimental Procedure

Nanoparticles of GaP have been prepared by mild aqueous synthesis and 2 low temperature methods of colloidal chemistry [1, 9-12]. The capping agent - trioctilphosphine oxide (TOPO) was added to the reaction mixture to prevent coagulation of the GaP nanoparticles and as the result the particles with the 10-60 tirn dimensions have been obtained.

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The essential details of these methods as the base for comparison of nanoparticles prepared at different conditions are noted below. The used in mild aqueous synthesis NaOH pellets were dissolved in distilled water. Ga203, red or white phosphorus and h were mixed and added to the NaOH solution. The mixed solution was then placed into an autoclave and heated there in an oven for 8 hours at 200 or 125°C. After the completion of heating, accompanying by intense ultrasonication or without it, the autoclave was taken out of the oven and cooled. GaP nanoparticles also have been obtained taking advantage of the reaction of Ga(OH)4 with PH3 which was produced from white phosphorus dispersed in alkali solution:

P4 + 30H + 3H20 — PH3 + 3H2P02" Ga203 + 20H' + 3H20 — 2Ga(OH)„"

Ga(OH)4" + PH3 -» GaP + 3H20 +OH"

In order to improve the yield of GaP, iodine was added to induce the reaction with white phosphorus, based on follow process:

P4 + 2I2 +40H" +4H20 -» 2PH3 + 2H3P04 +41

In the colloidal method the GaP nanoparticles were synthesized through the reaction between anhydrous gallium chloride (GaCb) and sodium phosphide (Na3P). Na3P was used in an Ar-filled glove-box through the reaction of white P and Na in toluene. The reaction went for 8 hrs. at 110°C, ultrasonication and mechanical mixing until completion of the process with creation of black suspension of Na3P in toluene. For the syntheses of GaP nanoparticles 0.8 g of GaCb were dissolved in 30 mL of anhydrous toluene and the solution mixed under heat at 100°C. Then 1.5 g of Na3P was dissolved in 50 mL of toluene and added to the GaCb solution which was reacted for 2 hrs at 100°C under mechanical mixing. GaP nanoparticles were obtained as the result of the reactions:

12Na + P4 = 4Na3P GaCl3+Na3P—GaP+3NaCl

The obtained nanoparticles was filtered, washed with ethanol or other special solvents under ultrasonic and ultracentrifuge treatment for separation in dimensions and preparation of a suspension for any nanocomposite. Poly (2-vinylpyridine) (P2VP), biphenyl vinyl ether (BPVE) and tetrahydrofuran (THF) polymers were used for preparation on GaP nanocomposites. Thickness of the polymer composite film was within 250-300 nm as determined by Atomic Force Microscopy (AFM). Thermo-gravimetric analysis (TGA) was used to evaluate the decomposition of the nanocomposites. The next procedures have been used in fabrication of the nanocomposites [12]: 1) GaP powder was ultrasonicated in methylethylketone (MEK) using Branson 5210 ultrasonic bath. Then, PGMA was added to the MEK solution. GaP to polymer ratio was less than 1:10. 2) GaP powder was dispersed in water-ethanol mixture (1:1 volume ratio) and ultrasonicated using Branson 5210 bath for 120 min. Then, PGMA-co- POEGMA was added in the form of water-ethanol mixture (1:1 volume ratio) solution. GaP to polymer ratio was less than 1:3. Nanocomposite films were deposited on quartz slides via dip-coating;

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3) GaP powder was dispersed in the biphenyl vinyl ether/dichloromethane (BPVE/DCM) solution; the solution was stirred and filtered from the excess of the powder. A few mL drops of the settled solution were casted onto silicon wafer. Note, that the light emissive nanocomposites of the best quality can be obtained only from a suspension that is thoroughly washed, ultrasonicated and subsequently dried. The dried powders of GaP nanoparticles and GaP nanocomposites were then characterized using standard methods of Raman light scattering (RLS), X-ray diffraction (XRD) and photoluminescence (PL). For comparison, an industrial and specially grown and aged GaP [2-8, 13,13] single crystals were used. The instruments for Raman light scattering and luminescence included Spectrograph Triax 552 interfaced to Symphony CCD detection system and Coherent Innova 100-3K Krypton Ion Laser System. The spectra of Raman scattering was obtained at room temperature by excitation with 752.5 nm radiations and calibrated with the relevant étalons. X-ray diffraction data were collected on Rigaku ULTIMA IV powder diffractometer using a monochromator and Cu Ka radiation (1.5406 Â). All scans were in the 6-9 mode at 300 K. Luminescence (PL) was excited by the N2 laser nanosecond pulses at wavelength 337 nm and measured at room temperature.

Results and Discussion

Figure 1 shows spectra of Raman light scattering from GaP nanoparticles prepared by aqueous synthesis at different conditions. Characteristic GaP Raman lines from the doped perfect GaP single crystals as well as from the nanoparticles prepared at decreased temperature (125°C) using white P were narrow and intense (Figure 1, spectrum 1 and 4 respectively), whereas they were weak and broad from nanoparticles prepared using red P at increased (200°C) temperature (Figure 1, spectra 2 and 3) and especially from not thoroughly washed nanoparticles (Figure 1, spectrum 2).

Figure 1. RLS from GaP nanoparticles of different treatment prepared from white or red P (spectra 2-4) in comparison with perfect GaP bulk single crystals (spectrum 1). Please see the

details in the text.

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Figure 2. X-ray diffraction from GaP nanoparticles in comparison with the diffraction from perfect GaP single crystal (spectrum 4). 1. White P, the best low temperature synthesis, well-

treated powder. 2. White P, not the best performance and powder treatment. 3. Red phosphorus, the best result.

Considerable changes in quality of GaP nanoparticles due to improvement of their preparation are shown in Figure 2, where one can see the most characteristic for GaP the 6-20 range of X-ray diffraction from GaP nanoparticles prepared at different conditions (spectra 1-3) in comparison with the diffraction from perfect GaP single crystal (spectrum 4). The nanoparticles obtained by low temperature aqueous synthesis using white phosphorus develop clear characteristic lines like those obtained from perfect GaP bulk single crystals taken from our unique collection of long-term ordered GaP single crystals [8] (Fig. 2, spectra 1 and 4). Contrary to that, nanoparticles prepared on the base of red phosphorus and/or not under optimum conditions (temperature, ultrasonication) show broad, weak or no GaP characteristic lines (spectra 2 and 3). The halfwidth of the characteristic lines was used for evaluation of the average dimensions of nanoparticles equal to 10-100 nm dependently on the preparation conditions. Figure 3 compares the luminescence spectra of our long-term (up to 50 years) ordered GaP single crystals (spectrum 1) to that from high quality GaP nanoparticles [9, 11] and their GaP nanoparticles/polymers nanocomposites [10, 12]. The best quality GaP nanoparticles have been prepared by hydrothermal or colloidal synthesis from white phosphorus at decreased temperature (125°C) and intense ultrasonication. Nanocrystals stored as dry powder demonstrate rather broad luminescent band with maximum at 2.8 eV (spectrum 2), while the nanocrystals of about 10 nm sizes, thoroughly separated and distributed in a suspension, that prevent their coagulation, mechanical and optical interaction, exhibit bright narrow-band luminescence with maximum at 3.2 eV, approximately 1 eV above

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the position of the absorption edge in GaP at 300°K (spectrum 3). The thoroughly washed, ultrasonicated and dried nanopowders as well as their specially prepared suspensions have been used for fabrication of blue light emissive GaP nanocomposites on the base of some optically and mechanically compatible with GaP polymers. Comparing the results for the nanocomposites prepared from GaP powder or suspension (Fig. 3, spectra 2 and 3 respectively), it was established that the maximum shift to ultraviolet and the best quality in general have the nanocomposites obtained from the nanoparticles stored as a suspension in a suitable liquid.

Fig. 3. Luminescence of perfect bulk GaP single crystals (1) in comparison with the luminescence of GaP nanoparticles and GaP/polymers nanocomposites (2-3). Nanoparticles have been prepared using white P by mild aqueous or colloidal synthesis at decreased temperature and

stored as the dry powder (spectrum 2) or suspension in a liquid (spectrum 3).

According to our measurements, the matrix polymers PGMA-co-POEGMA or BPVE used in this work provide no contribution to the spectra of luminescence of the based on these matrixes, so, the nanocomposite spectra coincide with those obtained from the relevant GaP powders or suspensions. We note that in the GaP/BPVE nanocomposite the position of the luminescent maximum can be changed between 2.5 - 3.2 eV and the brightness is 20-30 more than in the PGMA and PGMA-co-POEGMA matrixes. We explain the broadening of the luminescence band and the shift of its maximum to low photon energies in luminescence of the nanocomposite based on the GaP powder by presence in the powder of the nanoparticles with the different dimensions between 10-100 nm. Meanwhile, the nanocomposites on the base of the suspensions containing only app. the 10 nm nanoparticles, exhibit bright luminescence with maximum at 3.2 eV due to high transparency of 10 nm nanoparticles for these high energy emitted photons and pronounced quantum confinement effects since this diameter equals the Bohr diameter of the bound exciton in GaP. In accord with our data [9-14] the shift due to the quantum confinement effects is about a few tenths of eV and, obviously, it is impossible to explain only through this effect the dramatic 1 eV expansion of the region of luminescence at 300 K to the high-energy side of the spectrum. In order to explain this interesting phenomenon we postulate that the nanocrystals, much like the ideal long-term ordered bulk GaP single crystals, exhibit this huge increase in blue-shifted luminescence due to: (a) negligibly small influence of defects and non-radiative recombination of electron-hole pairs and very high efficiency of their radiative annihilation, (b) high perfection

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of nanocrystal lattice, and (d) high transparency of nanocrystals due to their small dimensions for the light emitted from high points of the GaP Brillouin zones, for instance, in the direct transitions IV - Ti5

v between the conductive and valence bands with the photon energy at 300K equal to 2.8 eV [15] and (e) high efficiency of this so called "hot" luminescence.

Conclusions

First attempts to prepare GaP nanoparticles [1] yielded room temperature luminescence with maximum shifted only to 2.4 eV in comparison with the new maximum at 3.2 eV. It confirms significant achievements in technology of QaP nanoparticles and GaP/polymers nanocomposites. On the base of these improved technologies for preparation of GaP nanoparticles and GaP/polymer nanocomposites we can change within the broad limits the main parameters of luminescence and expect to create a framework for novel light emissive device structures using dramatic 1 eV expansion of GaP luminescence to UV region.

Acknowledgements

The authors are very grate&l to the US Dept. of State, Institute of International Exchange, Washington, DC, the US Air Force Office of Scientific Research, Science & Technology Center in Ukraine (STCU), Clemson University, SC, Istituto di elettronica dello stato solido, CNR, Rome, Italy, Universita degli studi, Cagliari, Italy, Joffe Physico-Technical Institute, St. Petersburg State Technical University, Russia and Academy of Sciences of Moldova for support and attention to our extended (1963-2009) research efforts. Our special gratitude to Dr. E. Rusu from Institute of Electronic Engineering and Industrial Technologies, post graduate students S. Bilevschii and A. Racu from Institute of Applied Physics, Moldova, Prof. I. Luzinov, Prof. G. Chumanov, Dr. B. Zdyrko and Dr. D. Van DerVeer from Clemson University, SC, for their contribution in elaboration of GaP nanotechnology, characterization of nanoparticles and nanocomposites.

References

1. S.L. Pyshkin, J. Ballato, G. Chumanov, J. DiMaio, and A.K. Saha, Preparation and Characterization of Nanocrystalline GaP, Technical Proceedings of the 2006 NSTI Nanotech Conference, 3, 194-197 (2006). 2. S.L. Pyshkin, R.P. Zhitaru, and J. Ballato, Modification of Crystal Lattice by Impurity Ordering in GaP, Int. Symposium on Defects, Proceedings of the MS&T2007 Conference, International Symposium on Defects, Transport and Related Phenomena (Detroit, MI, September 16-20, 2007), 303-310 (2007). 3. S.L. Pyshkin, J. Ballato, M. Bass, and G. Turn (invited), Luminescence of Long-Term Ordered Pure and Doped Gallium Phosphide, TMS Annual Meeting, Symposium: Advances in Semiconductor, Electro Optic and Radio Frequency Materials (March 9-13, New Orleans, LA). J. Electronic Materials, 37(4), 388-395 (2008). 4. S. Pyshkin, and J. Ballato, Long-Term Ordered Crystals and Their Multi-Layered Film Analogues, Proceedings of the 2008 MS&T Conference, Pittsburgh Symposium on Fundamentals & Characterization, Session "Recent Advances in Growth of Thin Film Materials", 889-900 (2008). 5. S. Pyshkin, R. Zhitaru, J. Ballato, G. Chumanov, and M. Bass, Structural Characterization of Long Term Ordered Semiconductors , Proceedings of the 2009 MS&T Conference, International Symposium "Fundamentals & Characterization," 698-709 (2009).

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6. S. Pyshkin, J. Ballato, M. Bass, G. Chumanov, and G. Turn, Properties of the Long-term Ordered Semiconductors, The 2009 TMS Annual Meeting and Exhibition, Suppl. Proc, (San Francisco, February 15-19, 2009), 3,477-484 (2009). 7. S. Pyshkin, J. Ballato, M. Bass, and G. Turri, Evolution of Luminescence from Doped Gallium Phosphide over 40 Years, J. Electronic Materials, 38(5), 640-646 (2009). 8. Sergei Pyshkin and John Ballato, Evolution of Optical and Mechanical Properties of Semiconductors over 40 Years, J. Electronic Materials, Springer, DOI: 10.1007/sl 1664-010-1170-z, 39(6), 635-641 (2010). 9. S. Pyshkin, J. Ballato, G. Chumanov, N. Tsyntsaru, E. Rusu, Preparation and Characterization of Nanocrystalline GaP for Advanced Light Emissive Device Structures, The 2010 Nanotech Conference (Anaheim, CA, June 21-24), NSTI, NSTI-Nanotech 2010, www.nsti.org, ISBN 978-1-4398-3401-5,1, 522-525 (2010). 10. S. Pyshkin, J. Ballato, I. Luzinov, B. Zdyrko (2010) Fabrication and Characterization of the GaP/Polymer Nanocomposites for Advanced Light Emissive Device Structures, The 2010 Nanotech Conference (Anaheim, CA, June 21-24), NSTI, NSTI-Nanotech 2010, www.nsti.org, ISBN 978-1-4398-3401-5, 1, 772-775 (2010). 11. Pyshkin SL, Ballato J, Belevschii S, Rusu E, Racu A, Van DerVeer D, Synthesis and Characterization of GaP Nanoparticles for Light Emissive Devices. The 2011 Nanotech Conference (Boston, MA, June 13-16), NSTI, NSTI-Nanotech 2011, www.nsti.org, ISBN 978-1 4398-7142-3 1, 327-330 (2011). 12. Pyshkin S. L., Ballato J., Luzinov I. & Zdyrko B., Fabrication and characterization of GaP/polymer nanocomposites for advanced light emissive device structures, Journal of Nanoparticle Research (2011). DOI 10.1007/sl 1051-011-0547-0 13. Sergei L. Pyshkin and John M. Ballato, Long-Term Convergence of Bulk- and Nano-Crystal Properties, Advances and Applications in Electroceramics: Ceramic Transactions, 226, 77-90 (2011). 14. Sergei L. Pyshkin and John Ballato, Long-Term Convergence of Bulk- and Nano-Crystal Properties, Chapter 19 in "Optoelectronics - Materials and Technics", ISBN 978-953-307-276 0, InTech - Open Access Publisher, 459-475 (2011) 15. Zallen R. & Paul W., Band Structure of Gallium Phosphide from Optical Experiments

at High Pressure, Phys. Rev. 134 A1628-A1641 (1964).

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POSITRON LIFETIME ANALYSIS OF POLYUREA-NANOCLAY COMPOSITS

Naidu V. Seetala1, Gabriel Burks1, Danny Hubbard2, Alex Trochez2 and Valéry N. Khabashesku3

'Department of Physics, Grambling State University, Grambling, LA 71245, naidusv@gram.edu department of Chemistry, Grambling State University, Grambling, LA

department of Chemical and Biomolecular Engineering, University of Houston, Houston, TX

Keywords: Aliphatic Polyurea, Nanoclay, Positron Lifetime, Positronium, Pore Structure

Abstract

Positron lifetime spectroscopy (PLS) is used to study 1 -5 wt% nanoclay incorporated aliphatic polyurea films prepared by two different methods. Set-1 used aliphatic polyaspartate polyurea system consists of aliphatic diisocyanate resin, while set-2 used pre-made aliphatic-polyurea granules. PLS was used to study pore structure in polyurea samples. The third lifetime component, related to positronium formation in free spaces, provided the information on pore size and concentration of pores. The third lifetime component showed ~ 2 ns for set-1 and ~ 1.8 ns for set-2 with relative intensities of 16% and 19%, respectively; indicating that set-1 polyurea has larger pores with lesser concentration compared to set-2 polyurea. There is good correlation between positronium lifetime parameters (both lifetime and intensity) and % nanoclay in set-1 polyurea, but no correlation is observed for set-2 polyurea. The set-2 films showed bad quality and the polymer curing is not as good as set-1 films.

Introduction

Nanoclay reinforced polymer composites improve multifunctionalities such as structural and thermal properties [1, 2]. Polyurea with nanoclay is important in developing appropriate composite material for the blast and fragmentation resistant applications. Understanding the mechanisms involved in improving the mechanical strength of polyurea with nanoclay is important in developing appropriate composite material for the blast and fragmentation resistant applications. Free volume parameters, such as pore size and pore concentration in polymers, play an important role in determining the physical properties of the polymer. Positron Annihilation Techniques (PAT) are very sensitive to the microporosity of the sample [3, 4], Positron lifetime spectroscopy (PLS) is very sensitive to free volume in polymers. The third lifetime component, related to positronium formation in free spaces in polymers, provides the information on pore size (proportional to positronium lifetime) and concentration of pores (proportional to the intensity of the positronium component). Here, we use PLS to study the pore structure variations in aliphatic polyurea with 1-5 wt% nanoclay prepared by two different procedures: 1) starting with aliphatic polyaspartate polyurea system consists of aliphatic diisocyanate resin (set-1), and 2) starting with pre-made aliphatic-polyurea granules (set-2).

Experimental

Two sets of aliphatic polyurea (PU) samples were prepared with 1-5 wt% nanoclay. The first set of organoclay-polyurea composites have been prepared [5] from castable aliphatic polyaspartate polyurea system acquired from Bayer Materials Research Company. The system consists of aliphatic diisocyanate resin (component A). Commercial nanoclay product, Cloisite 15A,

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functionalized with a quaternary amine containing a C14-C18 long-chain alkyl tails, was obtained from Southern Clay Products. The nanoclay powder was exfoliated by placing it into a mixture of secondary polyetheramines (Component B) and tiien stirred at 80 °C for 3 hrs. This was followed by degassing of the dispersion and of aliphatic resin in vacuum oven at 75 °C. The components were mixed in desired ratio to get 1-5 wt% nanoclay by placing the PU components A and B into a dual cartridge gun equipped with a mixing pipe dispenser purchased from PLUS PAK, then cast into a Teflon mould and left for curing overnight. A second set of aliphatic polyurea-nanoclay composites has been prepared starting with gently warming the pre-made polyurea obtained from Versaflex to 37 °C, added 1-5 wt% montmorillonite nanoclay, heated the mixture to 75 °C and stirred for 2 hours, placed the mixture into a ceramic plate containing a 2 cm well to make a mold, and allowed to dwell for 24 hours in a convectional oven at 90 °C followed by letting it stand for 1 hour in vacuum oven.

Positron lifetime system is used to study porous structure in 1-5 wt% nanoclay incorporated both sets of aliphatic polyurea films. A 22Na source is sandwiched between two identical polyurea films under study and the lifetime spectrum was collected. When a positron is emitted from the source, a high energy (1.28 MeV) nuclear gamma is simultaneously emitted from the source, which serves as birth signal and detected by the first scintillator. The positron (antiparticle of electron), which enters the sample, is trapped in a lattice defect and finally annihilates with an electron around the defect site and gives two low energy (0.511 MeV) annihilation gammas, which serve as the death signal and detected by the second scintillator. The time delay between the birth and deam signals gives the positron lifetime at the annihilation site in the sample. The positron lifetime depends on the average electron density around the defect site, which is a function of the defect size and type. As the defect size increases the electron density at the annihilation site decreases, thus the corresponding positron lifetime increases. A 16 ns delay is introduced for the time calibration of the spectrometer and found to be 0.0125 ns/ch. The positron lifetime spectra were analyzed using POSFIT computer program [6] and obtained three lifetime components for all samples. The first and second lifetime components are related to positron annihilation with electrons with in polymer chains and between polymer chains, respectively, without forming a positronium (hydrogen like) atom. As the electron concentration at the annihilation site increases the corresponding positron lifetime increases. The electron concentration within the free volume (pores) of the polymer is so low such that a positron can find an isolated electron to combine to form a positronium atom that eventually decays to gamma rays. The third lifetime component is related to the positronium lifetime in these films, which is related to free volume in polymer. The positronium lifetime is proportional to the pore size. The relative intensities of these three lifetime components are proportional to the number of positrons annihilated in these three states described which are directly related to the concentrations of these different sites available within the polymer.

Results and Discussion

Figure 1 show the positron lifetime spectra for two sets of blank aliphatic polyurea films prepared by different methods described above using 22Na positron source along with a time resolution peak collected using 60Co source ( Co emits two high energy gammas simultaneously, thus it serves a instrumental resolution function to use in the POSFIT program). The third lifetime component has longer lifetime in set-1 sample compared to set-2 sample. The difference can be clearly seen at the higher time scale (right edge of the lifetime spectrum). Indicating the set-1 sample has larger pores compared to set-2 sample. More detailed analysis is described below.

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Figure 1: Positron lifetime spectra for two sets of aliphatic-polyurea prepared by different methods along with time resolution function.

The positron lifetime results for Set-1 aliphatic polyurea with 1-5 wt% nanoclay samples are summarized in Figures 2 and 3, where the third lifetime positronium component variations are shown with varying wt% nanoclay in the aliphatic polyurea. The positronium lifetime decreases with the wt% nanoclay in the polyurea, indicating that the pore size in the polymer decreases with the percentage nanoclay in the polymer. The positronium intensity increases with the wt% nanoclay in the polyurea, indicating that the pore concentration in the polymer increases with the percentage nanoclay in the polymer. This may indicate that the larger pores are breaking into smaller pores as the percent of nanoclay increases in the polymer.

Figure 2: Positron lifetime third (positronium) component with % nanoclay in polyurea (set-1)

Set-1 samples are most consistent and reproducible compared to the set-2. The set-2 films are of bad quality and the curing of the polymer film is not good. Thus the positron lifetime results for set-2 did not show systematic variations with % nanoclay as in the case of set-1. So, we will only make a qualitative comparison between set-1 and set-2 samples, and the curing process in the process for set-2 samples need to be improved in order to obtain good quality samples.

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The third component lifetime (T3) and its relative intensity (I3) for blank polyurea prepared in set-2 have values of 1.8 ns and 19%, respectively. These values for set-1 samples showed ~ 2.0 ns lifetime with 16% intensity. These results indicate that set-1 polyurea has larger pores with lesser concentration compared to set-2 polyurea. There is no good correlation between positron lifetime components (both T3 and I3) and % nanoclay in set-2 polyurea, unlike set-1 polyurea. We are trying to improve the consistency in the polymer film preparation for set-2 procedure and investigating to improve the curing step in set-2 preparation methods. Set-1 polyurea preparation procedure seems to be better compared to the set-2 preparation procedure.

Figure 3: Intensity of the positron third lifetime component with % nanoclay in polyurea (set-1).

Conclusions

Aliphatic polyurea films with 1-5 wt% nanoclay are prepared by two different methods (set-1 and set-2). Set-1 used aliphatic polyaspartate polyurea system consists of aliphatic diisocyanate resin, while set-2 uses pre-made aliphatic-polyurea granules. Positron lifetime spectrometer is used to characterize the pore structure (relative pore size and pore concentration) using the positronium (third) lifetime component. The positronium lifetime decreases with the % nanoclay in the set-1 aliphatic polyurea as the % nanoclay increased, indicating a decrease in pore size with increasing % nanoclay. The positronium intensity increases with the % nanoclay, indicating that the pore concentration increases with % nanoclay in the set-1 polymer. The set-2 films, prepared from pre-made polyurea granules, showed bad quality, less consistency, low reproducibility, and the curing of the polymer film is not as good as set-1 films. The positron lifetime results indicate that set-1 polyurea has larger pores with lesser concentration compared to set-2 polyurea. There is a need to improve the curing process in the set-2 polyurea films preparation methods.

Acknowledgements

The work has been supported by a grant from Clarkson Aerospace, Air Force contract FA8650-05-D-1912 from the Air Force Research Laboratory Materials and Manufacturing Directorate.

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References

[1] H. H. Wang and K. V. Chen, "A novel synthesis of reactive nano-clay polyurethane and its physical and dyeing properties", J. Appl. Polym. Sei., 2007, Vol. 105, ppl581-1590.

[2] N. Sheng, M. C. Boyce, D. M. Parks, G. C. Rutledge, J. J. Abes, and R. E. Cohen, "Multiscale Micromechanical Modeling of Polymer/Clay Na-nocomposites and the Effective Clay Particle", Polymer, 2004, Vol. 45(2), pp.487-506.

[3] Y. Itoha, A. Shimazu, Y. Sadzuka, T. Sonobe and S. Itai, "Novel method for stratum corneum pore size determination using positron annihilation lifetime spectroscopy", International J. Pharmaceutics, 2008, Vol. 358, p-91-95.

[4] S. V. Naidu, S. V. Sastry, C. S. Maxie and M. A. Khan, "Effect of Plasticizer on Free Volume and Permeability in Cellulose Acetate Pseudolatex Membranes Studied by Positron Annihilation and Tracer Diffusion Methods", Mat. Sei. Forum, 1997, Vol. 255-257, pp333-335.

[5] N. Seetala, G. Burks, D. Hubbard, A. Trochez, and V. Khabashesku, ACS Division of Polymeric Materials: Science and Engineering Preprint 2011.

[6] P. Kirkegaard, N. J. Pedersen, and M. Eldrup, a program from Riso National Laboratory (1989).

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Rheological properties of suspensions of nanopowders Yttr ium Oxide (V2O3) and Magnesium-Aluminum spinel (MgAl2Oi)

G. Zyla1, M. Cholewa1, A. Witek2, J.P. Plog3, V. Lehmann4, T. Oerther4, D. Gross4

1 Department of Physics, Rzeszow University of Technology, Aleja Powstancow Warszawy 6, 35-905 Rzeszow, Poland,

2 Institute of Ceramics and Building Materials, Warsaw, Poland, 3 ThermoFisher Scientific, Karlsruhe, Germany,

4 Bruker Biospin GmbH, Rheinstetten, Germany.

Keywords: Rheology, Rheo-NMR, Viscosity, Nanopowder, Nanoparticiples, Nanopowder suspension, Ceramics

Abstract

We report results of rheological experiments on suspensions of nanopowders. We use ce-ramic nanopowders such as Yttrium-Aluminum garnet (YAG - Baikowski), Magnesium-Aluminum spinel (S30CR Baikowski - MgAl204) and Yttrium Oxide {Y203). These nanopowders have been suspended in alcohols. Rheological studies have been carried out using Haake Mars 2 rheometer. The samples were tested and then compacted (sub-jected to centrifugation) in a centrifuge and sintered. We have been studying the link of the rheological properties of suspensions of nanopowders at different concentrations with the mechanical properties of the final produced ceramics.

Introduction

Rheology is the study of the flow of matter, primarily in the liquid state, but also as 'soft solids' or solids under conditions in which they respond with plastic flow rather than deforming elastically in response to an applied force [1]. It applies to substances which have a complex molecular structure, such as muds, sludges, suspensions, polymers and other glass formers (e.g. silicates), as well as many foods and additives, bodily fluids (e.g. blood) and other biological materials.

Newtonian fluids can be characterized by a single coefficient of viscosity for a specific temperature. Although this viscosity will change with temperature, it does not change with the flow rate or strain rate. Only a small group of fluids exhibit such constant viscos-ity, and they are known as Newtonian fluids. But for a large class of fluids, the viscosity change with the strain rate (or relative velocity of flow) and are called non-Newtonian fluids. Rheology generally accounts for the behavior of non-Newtonian fluids, by charac-terizing the minimum number of functions that are needed to relate stresses to the rate of change of strains or strain rates. Similar developments has been performed worldwide [1],[2],[3]. The research, performed by our team, has produced new and practical results which could be used to produce nanomaterials with required properties in the future. The use of Rheo-NMR [4],[5],[6],[7],[8] and RheoScope [9] has added a new dimension to our project.

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Much of nanomaterials (such as nanopowders Y203 and MgA^O^) has not been suf-ficiently well studied rheologically. Our group in RUT has undertaken research on the rheological properties of suspensions of ceramic nanopowders in ethanol. This is impor-tant for industry because the suspension is subjected to further technological process lead in the manufacture of these ceramics. Some material was also tested in research laboratories and companies e.g. ThermoFisher Scientific and Bruker Biospin GmbH in Germany. This work presents a summary of the results obtained by us for nanopowders.

Experimental

The materials that we measure was Yttrium Oxide (Y2O3) and Magnesium-Aluminum spinel (MgA^Oi) received from ICBM in Warsaw, Poland.

Preparation of the samples in ICBM.

The sample of V2O3 ceramic nanopowder was purchased from a commercial company and modified at ICBM by grinding to smaller grains. The second sample used in our ex-periments has been commercially available nanopowder of Magnesium-Aluminum spinel (MgAl204) manufactured by Baikowski. Several tests have been performed on both samples including (a) imaging with an electron microscope, (b) measurements of the size of grains with X-Ray Diffraction (XRD) and (c) elemental analysis with X-Ray Fluores-cence (XRF). The average size of nanoparticles measured with XRD was (31.7±0.4) nm for Y2O3 and (40 ± 1) nm for Magnesium-Aluminum spinel. Images taken with SEM mi-croscope clearly shows tendency for these nanomaterials to create larger conglomerates.

Preparation of the sample for rheological measurment.

The tests were performed on samples containing different concentrations of nanopowders in alcohol. Samples were prepared on an analytical balance Radwag WAS 220/X with the order of accuracy of 0.1 mg. The first phase of preparing the material for the study was by placing the nanopowder in a glass vessel and accurate measurement of the required concentration. Then the vessel was filled with an appropriate amount of ethyl alcohol (96% pure p.a. produced by POCH, CAS: 64-17-5). The sample was subjected to mechanical stirring for 30 minutes and then placed it in an ultrasound wave bath (Ultron U-505) for 60 minutes. Subjecting the material to the ultrasound resulted in the breakdown of agglomerates of nanopowders.

Rheological measurements at RUT.

Study of rheological properties of suspensions of nanopowders were made at the RUT on Haake Mars 2 rheometer from Thermo Electron Corporation in the geometry of the measuring a double cone (diameter 60 mm, cone angle 1°). To monitor the temperature, a Peltier system and thermostat Phoenix 2 from Thermo Electron Corporation was used. In addition, measuring system was isolated from the environment by using glass rings. Performed measure dynamic viscosity in the range of shear rates from 0,01 s_1 to 2000 s_1

and the temperature range from —15 °C to 20 °C.

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Experiment in ThermoFisher Scientific.

In the laboratory of ThermoFisher Scientific in Karlsruhe, Germany, rheological studies were performed on a rheometer Haake Mars 3 equipped with RheoScope module which allows the optical observation of samples during the rheological tests. To perform the measurements used in measuring the plate-plate geometry with a diameter of 60 mm and carried out research in the range of shear rates 1 s_1 to 1000 s - 1 on samples of 1% and 5% concentration of Y203.

Measurment in Bruker Biospin.

To make the experiments in a laboratory Bruker BioSpin in Rheinstetten, Germany a Rheo-NMR system was used. The measuring system consisted of AV 300 WB III equipped with a Rheo-NMR module. Measuring geometry of a plate-cone (diameter 16 mm, cone angle 7°) was used and studied the flow velocity distribution inside the sample at speeds 2.5 rad/s and 5.0 rad/s for a concentration of 10% Y2O3 nanopowder. With NMR it was also possible to perform testing of the pure ethyl alcohol obtained directly from the original manufacturers pack before and after subjecting it to ultrasound.

Results and Discussion

As a result of our measurements we observed differences between the dynamic viscosity of ethyl alcohol tested immediately after downloading it from the packaging manufacturer and the dynamic viscosity of alcohol after subjecting it to ultrasound in ultrasound wave bath. The results of this observation is shown in Figure 1. In the laboratory of

Figure 1: Summary of dynamic viscosity of ethyl alcohol before being subjected to the action of ultrasound and after. Both studies were carried out at —15 °C.

Bruker Biospin NMR study was conducted on both samples. The study showed that the chemical composition of alcohol does not change. We assume that the viscosity change

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is related with the removal of "airbubbles" from the volume of alcohol, which affects the viscosity, but not on the chemical composition.

Our laboratory at RUT is air-conditioned from 7 a.m. until 4 p.m.. We have observed that the time of day in which the prepared samples has a significant influence on the measurement results. The results of measurements of samples prepared in the morning are different from the results of measurements of samples prepared in the afternoon. This dependence is shown in Figure 2. We believe that periodic changes in temperature inside the laboratory affect the operation of the equipment used for sample preparation, in particular ultrasound wave bath. As shown in Figure 1 ultrasound have a significant effect on the viscosity of the sample.

Figure 2: Comparison of viscosity of a suspension of 3% nanopowder Y2O3 in ethyl alcohol at — 5 °C.

Yttrium Oxide (Y203) Rheological studies in RUT were performed in CR mode (control rate) in the range of shear rates from 0,01 s_1 to 2000 s_1. Studies have shown that for low shear rate dynamic viscosity of the suspension grows to a certain limiting value of shear rate and then starts to decrease. Figure 3 shows the comparison of the dynamic viscosity of a suspension of 10% Y203 in alcohol at different temperatures. Measurements made at ThermoFisher Scientific enabled us to optically observe the interior of the sample during rheological testing we observed sedimentation of nanopowders. We observed that there exist limiting of shear rate at which nanoparticiples are "picked up" from the bottom of the measuring cell. For the measurement geometry used in the measurement (plate-plate 60 mm) it was about 120 s_1 for concentration of 1%. For the concentration of 5% it was at about 530 s_1.

Measurements made in the RUT were confirmed in the Laboratory Bruker Biospin, Germany. Measurement were performed for a plate-cone (diameter 16 mm, cone angle

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Figure 3: Results of dynamic viscosity measurments of suspension of 10% Y203 in ethylen in various temperature.

7°) measuring geometry. Performed two series of measurement at speeds 2.5 rad/s and 5.0 rad/s for a concentration of 10% Y2O3 nanopowder. Figure 4 shows the velocity distribution inside the sample, depending on the distance from the stationary plate. In the case of Newtonian fluid distribution of velocity between the stationary plate and rotating rotor is linear. As can be seen in Figure 4 velocity distribution in the test sample is non-linear. The graph shows the velocity distribution in the sample at the place marked on the image vertical line. At 2.5 rad/s velocity distribution is linear.

Also made measurements dependence of viscosity on the concentration of nanopowder in suspension. We noticed that the dynamic viscosity increases with increasing concen-tration nanopowder, which is presented in Figure 5.

Magnesium-Aluminum spinel {MgAl20^j Research viscosity of the suspension nanopow-der MgA\20\ in ethanol in the RUT was performed under the same environmental con-ditions as research of suspensions of Y2O3. The results are decidedly different, depending on the type of nanopowder used in the study. In the case of Y203 we had to deal with sedimentation and a critical value of shear rate, followed to a " pick up" nanoparticiples from the bottom of the measuring geometry. In the case MgAl2Ox nanopowder did not observe this phenomenon. In Figure 6 summarizes the results of measurements of dynamic viscosity MgAl20^ depending on temperature. It is clear that temperature does not have such a significant effect on the viscosity of the suspension as it was in the case of a suspension of Y203. We measured the viscosity dependence on concentration MgAl2Oi, as shown in Figure 7. From the measurements it is clear that the concen-tration has a significant effect on the viscosity of suspension. At lower concentrations (1%, 5%) after an initial nonlinear dependence of viscosity on the shear rate is linear, we

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Figure 4: Distribution of velocity in 10% Y203 in ethylen at rotor speed of 5.0 rad/s.

Figure 5: Dependence of dynamic viscosity of the shear rate for various concentrations Y2O3 in ethyl alcohol at 0 °C.

can therefore determine that the suspension as a Newtonian fluid. In the case of higher concentration (15%, 20%) observed a decrease of viscosity with increasing shear rate in the entire measuring range.

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Figure 6: Dependence of dynamic viscosity of 10% concentration of MqAliO^ in ethyl alcohol for various temperatures.

Figure 7: Dependence of dynamic viscosity of MgAhOï in ethyl alcohol on concentration at 0 "C.

Conclusions

Due to the enormous impact the environment during sample preparation, studies of The-ological properties of nanomaterials is difficult and requires great precision. We observed

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differences between the properties of samples prepared at various states of enviroment. Our research has shown that Theological properties of suspensions of different nanopow-ders have significantly different properties. We noticed that for suspensions of nanopow-der Y2O3 the sedimentation is almost immediate. To prevent falling nanoparticples the usage higher share rates is necessary. This is important for industry because it allows one to specify the minimal shear rate used in technological processes. The study of suspension of MgAl2Ot in the ethylen alcohol showed that the temperature does not significantly affect its viscosity. However, the change of concentration has significant in-fluence on the behavior of the sample and leads to an increase of viscosity and modyfies the material from a Newtonian one to non-Newtonian.

Acknowledgements

We would like to thank RHL-Service company from Poznan, Poland, in particular, Mr. Bartlomiej Borek, for arranging our visit in ThermoFisher Scietific Laboratory in Karl-sruhe, Germany.

References

[1] P. Oswald, Rheophysics - The Deformation and Flow of Matter, Cambridge Univer-sity Press (2009)

[2] G. Schramm, A Practical Approach to Rheology and Rheometry, Gebrueder HAAKE GmgH, (1994)

[3] E. Tamjid, Bernd H. Guenther, Rheology and colloidal structure of silver nanopar-ticles in diethylene glycol, Powder Technology, 197, 49-53 (2010)

[4] B.S. Douglass, R.H. Colby, L.A. Madsen, et al, Rheo-NMR of wormlike micelles formed from nonionic pluronic surfactants, Macromolecules, 41, 804-814, (2008)

[5] K. Ohgo, F. Bagusat, T. Asakura, et al, Investigation of structural transition of regenerated silk fibroin aqueous solution by Rheo-NMR Spectroscopy, Journal of the American Chemical Society, 130, 4182-4186, (2008)

[6] G. Mazzanti, E.M. Mudge, E.Y. Anom, In situ Rheo-NMR measurements of solid fat content, Journal of the American Chemical Society, 85, 405-412, (2008)

[7] P.T. Callaghan, Rheo-NMR and shear banding, Rheologica Acta, 47, 243-255, (2008)

[8] P.T. Callaghan, Elmar Fischer, Rheo-NMR: a New Application for NMR Microscopy and NMR Spectroscopy Bruker Report 149/2001, Institute of Fundamental Sciences-Physics, Massey University, Palmerston North, New Zeeland

[9] http://www.thermoscientific.com/ecomm/servlet/productsdetaiLl 1152-L11194-82133.13281391.-1

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Thermal Properties of Hemp-High Density Polyethylene composites: Effect of two different chemical treatments

Na Lu1, Shubhashini Oza2

'Department of Engineering Technology; University of North Carolina at Charlotte (UNCC), Charlotte, NC, 28223, USA

department of Civil & Environmental Engineering, UNCC, Charlotte, NC, 28223, USA

Keywords: Natural fiber composites, thermal properties, hemp fiber, surface modification

Abstract

The drive towards environmentally sustainable materials has resulted in use of natural fibers as an alternative to glass fibers for reinforcement in composites. Hemp has a greatest potential of being used as a reinforcing fiber since it is one of the strongest and stiffest natural fibers along with being biodegradable, low cost and fast growth. However, the incompatibility of hemp fiber with thermoplastic matrix results in poor interface adhesion and lower thermal and mechanical properties. In this study, we have systematically investigated the effects of two chemical treatments of hemp fibers on the thermal stability of hemp-HDPE composites. Composites of various volume fractions were synthesized with sodium hydroxide and silane treated hemp fibers. Thermo-gravimetric analysis was used for thermal stability characterization. The results indicate that chemical treatment of hemp fibers with silane treated yield composites with better thermal stability compared to that of composites treated with NaOH or untreated hemp-HDPE.

Introduction

Natural fiber reinforced composites are re-emerging as alternatives for synthetic fiber reinforced composites for the past two decades in automotive, building and construction industries[l]. The renewed interest in these composites are due to their unique material properties, including low cost, low density, non-abrasiveness, desirable mechanical and acoustical properties. Among the fibers processed from different parts of the plants, the bast fibers are the ones which are most frequently used in reinforcing composites. Some of the commonly used bast fibers are hemp, flax, kenaf, jute etc. Previous studies have shown that hemp fiber has excellent tensile strength of 814MPa (theoretical) with low density of 1.4g/cc, among the bast fibers[2]. Though these fibers have significant advantages, they have few drawbacks such as high moisture absorption and low thermal stability. To overcome these drawbacks, natural fibers are subjected to surface modifications such as physical or chemical treatment or use of coupling agents [3-5].

In this study, a systematic investigation on the effect of alkali treatment and silane treatment of hemp on the thermal properties of hemp-high density polyethylene (HDPE) composites has been studied. The polymer matrix used is HDPE since it is one of the polymers with very less branching and higher tensile strength.

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Experiment

Materials

Industrial hemp fibers with a length of 25 mm were obtained from Hempline Inc. (Ontario, Canada). The moisture content of the raw industrial hemp fiber ranged from 6% to 7%. HDPE pellets having a density of 0.953 g/cm3, a melt index (MI) of 0.38 g/10min, and a melting temperature range from 118°C to 131°C were obtained from Dow Inc. (U.S.A). NaOH pellets were obtained from Fisher Scientific (New Jersey, U.S.A), and the silane solution (triethoxyvinylsilane) was obtained from Sigma Aldrich Inc. (Missouri, U.S.A).

Alkali treatment

Hemp fibers were treated using 5wt% of NaOH solution. Fibers were immersed in the NaOH solution for 16 hours at 50°C followed by a wash with distilled water until their pH was neutral. Hemp fibers were then dried in the oven at 80°C for 10 hours and stored in desiccators prior to composite preparation.

Silane treatment

Hemp fibers were treated using 5wt% of silane solution. Silane solution was prepared by adding a 50v/50v ethanol/water mixture to the silane. 50v/50v acetic acid was added to adjust the pH to be in the range of 4-5. Fibers were then washed until their pH value was neutral. Hemp fibers were then dried in an oven at 80°C for 10 hours and stored in desiccators prior to composite preparation.

Composite preparation

In order to improve the surface adhesion of the natural fiber and thermoplastic matrix, a two-step manufacturing process was used. First, HDPE was processed separately into thin films using a C.W. Brabender 19.05mm single-screw extruder. Then each composite was manufacture by sandwiching a layer of fiber in between two layers of HDPE films, in a house made compression mold. Composites were prepared with 20%, 30%, 40% and 50% fiber volume fraction.

Thermogravimetric analysis

Thermal analysis of untreated and treated hemp fibers, pure HDPE, and composites were analyzed using a Mettler Toledo TGA. The analysis was conducted with a heating rate of 5°C/minute, from 25-600°C in a nitrogen atmosphere. For every sample five replicates were run.

Results and Discussion

The thermal stability of untreated hemp, along with NaOH treated and silane treated hemp is indicated in Figure 1. From this it is clearly seen that the degradation of hemp fiber, treated or untreated is a two-step process. The initial weight loss observed below 100°C was due to the removal of moisture. The drastic weight loss observed in the range of 245-330°C for untreated hemp could be attributed to the thermal degradation of

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hemicellulose, lignin and cellulose. But the same for treated fiber is between 250-375°C for NaOH treated hemp and 250-350°C for silane treated hemp.

Figure 1 : TGA curves for untreated, NaOH treated and silane treated hemp along with pure HDPE.

Both NaOH and silane modification increased the thermal stability of the hemp fiber in comparison to the untreated fiber. This could be attributed to reduction of the hemicellulose and lignin content after fiber treatment as observed in the FTIR spectra reported in our previous studies[6]. Among the two surface modifications studied, the thermal stability is slightly better with NaOH treatment than silane treatment.

The thermal degradation of the polymer is a one-step process (Figure 1). The degradation starts around 400°C and is complete by 475°C. Table 1 gives the details of degradation of untreated hemp, treated hemp and pure HDPE at various percentage weight losses.

Table 1: Degradation temperature (in °C) of untreated and treated hemp fibers

Pure HDPE Untreated 5wt% NaOH treated 5wt% Silane treated

Temperature at 10% weight loss (in °C)

428.18 265.15 305.76 291.91

Temperature at 15% weight

loss (in °C) 435.6 287.82 322.61 315.83

Temperature at 25% weight loss (in °C)

445.35 310.91 339.32 334.42

Temperature at 50% weight

loss (in °C) 459.10 336.84 357.62 351.09

The TGA curves for untreated hemp-HDPE composites at various fiber volume fractions along with pure HDPE are presented in Figure 2. From these curves we can see that the thermal stability of the composites is much lower than pure HDPE but higher than hemp fiber. Also, with increase in fiber volume fraction the thermal stability of the composite decreases.

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Figure 2: TGA curves for untreated hemp/HDPE composite at various fiber volume fractions.

Figure 3: TGA curves for NaOH treated hemp/HDPE composite at various fiber volume fractions.

Similar trends were observed for composites with NaOH treated hemp and silane treated hemp, which can be observed in Figure 3 and Figure 4, respectively. Similar results were obtained by Kim and et al, for rice husk flour reinforced polypropylene and HDPE composites[7]. Table 2 gives the details of degradation temperature for the all the three different composites studied at various fiber volume fractions.

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Table 2: Degradation temperature (in °C) of untreated and treated hemp fibers composites

Composite

20/80 Untreated hemp/HDPE 20/80 NaOH-hemp/HDPE 20/80 Silane-hemp/HDPE 30/70 Untreated hemp/HDPE 30/70 NaOH-hemp/HDPE 30/70 Silane-hemp/HDPE 40/60 Untreated hemp/HDPE 40/60 NaOH-hemp/HDPE 40/60 Silane-hemp/HDPE 50/50 Untreated hemp/HDPE 50/50 NaOH-hemp/HDPE 50/50 Silane-hemp/HDPE

Temperature at 10% weight loss (in °C)

338.5

343.17

361.83

326.25

334.25

344.25

320.00

333.67

340.08

317.50

328.67

328.17

Temperature at 15% weight loss

(in °C)

348.33

352.17

390.42

340.58

343.75

353.92

335.92

340.92

348.83

330.92

337.42

337.08

Temperature at 25% weight

loss (in °C)

388

400.83

421.67

357

382.75

399.08

351.42

353.92

383.42

343.50

348.25

347.17

Temperature at 50% weight loss (in °C)

450.75

450.83

452.00

441.67

450.33

450.00

415.58

441.42

448.83

404.25

418.42

421.33

By comparing the values in Table 1 and Table 2, it could be established that the thermal stability of the NaOH treated hemp is better than silane treated hemp. But for the composites, the trend is reversed. This could be explained due to the type of bond formation and changes that takes place due to silane and NaOH treatment on the hemp surface. During silane treatment, various type of reactions take place at the surface of the hemp fiber, such as hydrolysis, condensation, hydrogen bonding and covalent bond formation. Silanol molecules react with the hydroxyl group of the fiber resulting in formation of strong covalent bonds to the cell wall. Also the free silanols react with each other forming —Si-O-Si- bond. The vinyl group of the silane molecule couples with the thermoplastic matrix and increases the physical compatibility. Therefore, it could be concluded that the silane treatment not only modifies the surface of hemp, but when silane treated fibers are brought in contact with the polymer matrix, the organo-flinctional group of the silane molecule(in this case, vinyl) couples with the matrix and increases the strength of bonding. The enhanced covalent bonding added with physical compatibility with the polymer matrix increases the thermal stability of the silane treated hemp-HDPE composites. The NaOH treated hemp composites, neither result in covalent bonding on the fibers or increase physical compatibility with the matrix.

The thermal stability of as such hemp fiber, treated with NaOH is slightly higher than silane treated hemp fiber because NaOH is comparatively more effective in removing the

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hemicellulose and lignin content of the fiber. The details of the FTIR and SEM studies supporting these arguments are presented in the researchers' previous studies. [6, 8].

Figure 4: TGA curves for silane treated hemp/HDPE composite at various fiber volume fractions.

Conclusion

The thermal degradation behavior of untreated, NaOH treated and silane treated hemp fiber, along with their composites in HDPE matrix were investigated. The results obtained from the analysis of thermo-gravitation of above mentioned fibers and composites can be summarized as follows.

• Thermal stability of the treated fibers is better than untreated hemp fiber. Further, thermal stability of NaOH treated hemp is marginally better than that of silane treated hemp fiber. This could be explained due to the efficient removal of hemicellulose by NaOH treatment, resulting in more crystalline cellulose which increased the thermal stability of the fiber.

• Thermal stability of composites with treated hemp fiber is much better than the untreated hemp-HDPE composites. Both the treatments enhance the thermal stability of the composites. Though, thermal stability of the silane treated hemp-HDPE is much higher than the NaOH treated hemp-HDPE composites due to the formation of covalent bonding and physical compatibility of the organo-functional group of silane molecule with the matrix. NaOH treatment of hemp increased the number of-OH group on the surface thereby increasing reaction site for fiber/matrix interaction. These reaction sites resulted in hydrogen bonding and not strong covalent bonding as in the case of silane treatment.

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• The thermal stability of the both the untreated and treated hemp composites are in between that of the untreated hemp and pure HDPE. Also, with increase in fiber volume fraction, the thermal stability of the composites decreases, irrespective of untreated or treated hemp fiber.

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Reference

1. Li, X., L.G. Tabil, and S. Panigrahi, Chemical treatments of natural fiber for use in natural fiber-reinforced composites: A review. Journal of Polymers and the Environment, 2007. 15(1): p. 25-33.

2. Seber, E.H.L.a.D. Bast fiber application for composites. Available from: http://www,fibrealternatives.com/bast.fi ber.applacationforcomposite.htm.

3. Mukhopadhyay, S. and R. Fangueiro, Physical Modification of Natural Fibers and Thermoplastic Films for Composites - A Review. Journal of Thermoplastic Composite Materials, 2009. 22(2): p. 135-162.

4. Xie, Y.J., et al., Silane coupling agents used for natural fiber/polymer composites: A review. Composites Part a-Applied Science and Manufacturing, 2010. 41(7): p. 806-819.

5. Mohanty, A.K., M. Misra, and L.T. Drzal, Surface modifications of natural fibers and performance of the resulting biocomposites: An overview. Composite Interfaces, 2001. 8(5): p. 313-343.

6. Na Lu, S.M.B., Ian Ferguson Effect of Alkali and Silane Treatment on the Thermal Stability of Hemp Fibers as Reinforcement in Composite Structures, in The 2nd International Conference on Advances in Materials and Manufacturing Processes. 2011: Guilin, China.

7. Kim, H.S., et al., Thermogravimetric analysis of rice husk flour filled thermoplastic polymer composites. Journal of Thermal Analysis and Calorimetry, 2004. 76(2): p. 395-404.

8. Na Lu, S.M.B., Ian Ferguson The Effect of Alkylation on the Thermal and Mechanical Properties of Hemp Fiber Composite with Recycled High Density Polyethylene Matrix, in American Society for Composites 26th Technical Conference/Second Joint US-Canada Conference on Composites. 2011 : Montreal, Quebec, Canada

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DISCARDED ULTRAFINE PARTICIPATE CARBONACEOUS MATERIALS USED AS REINFORCERS OF RUBBER VULCANIZED PRODUCTS

'Guillermo R. Martin-Cortés; zFabio J. Esper; JLuiz Sâlvio Galvâo Dantas, 'Wildor T. Hennies, 2Francisco R. Valenzuela-Diaz

'Mining & Petroleum Eng. Dept. - Polytechnic School - University of Sao Paulo Av. Prof. Mello Moraes, 2373 Cid. Universitâria 05508-900 Säo Paulo, SP, Brazil,

2Raw Materials & Non-metallic Solids Laboratory - Metallurgie & Materials Eng. Dept. - Polytechnic School - University of Säo Paulo. Av. Prof. Mello Moraes, 2463 Cid. Universitâria 05508-900 Säo Paulo, SP, Brazil

3BENTONISA - Bentonita do Nordeste S. A. Joâo Pessoa, PB, Brazil Keywords: discarded ultrafine carbonaceous minerals (I), vulcanized rubber products (2), reinforcing raw materials for rubber products (3),

organic modified clays (4), nanocomposite material with organophillic clay (5)

Abstract

Vulcanized rubber products as tires, auto parts, sport-shoes, and others are widely spread. Raw materials for vulcanized rubber products come from industrial process generating C02 emissions. These products commonly use, at least 40phr of carbon-black filler obtained from petroleum in a contaminant process of incomplete burning of the fuel. NAOB® a rubber / organoclay nanocomposite material developed to make vulcanized rubber products without carbon-black is reinforced with only lOphr of organoclay. In order to keep the material specific volume, maintaining or increasing its technological characteristics, additional reinforcing materials are used. Between the additional materials, are ultrafine particulate materials discarded from mining of some carbonaceous minerals mat can be used directly from the mine or through an organic modification. This paper show some very good technological results obtained in rubber formulations after replacing the carbon-black by these set of materials.

Introduction

As known, rubbers vulcanized products are of black or a very dark color. Its technological performances as elongation; tensile properties; tensile strength; tensile stress, abrasiveness resistance, resilience, and impermeability among others depend in high ratio of the type of reinforcing filler added, traditionally, to the several carbon-black classes and their mixing proportions.

There are other reinforcing materials for rubber whose colors are not dark but clear other than black. In a general way, the most common non-black fillers for rubber are calcium carbonate, kaolin clay, precipitated silica, talc, barite, wollastonite, mica, precipitated silicates, fumed silica and diatomite. Of these, the three most widely used, by volume and by functionality, are calcium carbonate, kaolin clay and precipitated silica [1],

By the way, in the few last years, researchers from two departments of the Polytechnic School of the University of Sâo Paulo find that there are other type of reinforcing filler materials instead carbon-black and those non-black cited above. When these new reinforcing fillers materials are used the vulcanized rubber products present different appearance but also its technological performances are enhanced. As these materials are used in fewer quantities then carbon-black researchers seek complement materials to obtain same volumes of production. In this research line it was find that some nanoparticles materials from modified non-metallic minerals and also others nanoparticles substances from different sources, alone or mixed among them, used in

substitution of carbon-black as rubber reinforce filler are showing, an enhancing technological behavior in comparison with the traditional carbon-black or a similar behavior [2].

The characteristics which determine the properties a filler will impart to a rubber compound are particle size, particle surface area, particle surface activity and particle shape. Surface activity relates to the compatibility of the filler with a specific elastomer and the ability of the elastomer to adhere to the filler. Functional fillers transfer applied stress from the rubber matrix to the strong and stiff mineral. It seems reasonable then that this stress transfer will be better effected if the mineral particles are smaller, because greater surface is thereby exposed for a given mineral concentration. And if these particles are needle-like, fibrous or platy in shape, they will better intercept the stress propagation through the matrix. A compound's physical/mechanical properties can be strongly influenced by the surface activity of the filler, which is the ability of the filler's surface to bond with the matrix [1].

Carbon-black, material derived from oil is produced during the refining of crude oil or mineral oil with the known consequences of environmental contamination. Carbon-black physical properties are nanometric grain sized distribution and a depth black color. It is an amorphous carbon formed when the hydrocarbon is bumed without enough oxygen for it to burn completely. It's also known as lampblack, gas black, or channel black, is used to make inks, paints and rubber products. It can also be pressed into shapes and is used to form the cores of most dry cell batteries, among other things [3], The substitution of this substance, traditionally used as filler to reinforce the vulcanized rubber products, will help to reduce dependence on petroleum products and also to reduce environmental contamination.

NAOB® a rubber + organoclay nanocomposite [4] is a patented new type of material where smectite clay modified by ammonium quaternary [5] through low wet procedures and some additional materials, replaces carbon black as filler. Smectite clay has being used in several purposes, due to its excellent physical, chemistries, and structural properties. Main applications of smectites in Brazil are as agglomerative in métallurgie molds, in iron ore pelleting, as drilling mud in oil and water catch wells, as clarifier of wines, in catalyses at the chemistry and pharmaceutical industries, for pollutants removal in water treatment plants, in cosmetic among other applications.

In the composition of vulcanized rubber products is used, at least, 40 phr of carbon-black. It is not uncommon rubber compositions

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prepared for vulcanized with 50, 60 or more phr of carbon-black. As reinforce fillers NAOB <$> uses only 10 phr organoclay + 30 phr of additional materials, to achieve similar volumes of products to those obtained with carbon-black. Those additional materials could come from several different sources: from minerals with little or no industrial application, chemically modified or not, even different types of materials discarded by industry, also including those minerals discarded in the process of mineral extraction or end-discarded products as ashes from energy production [4].

This paper presents results and compare of: a) rubber without any reinforce material, b) traditionally vulcanized rubber products reinforced with carbon-black, c) vulcanized rubber products filled with organoclay and d) vulcanized rubber products were some minerals that have the element carbon in its chemical composition was used as reinforcing materials. Actually, some of these minerals are mining by-products without economical applications. Some of these minerals have not industrial applications because its physical properties are not adequate for the industries at the region around. Today destination of these minerals is the mine tailings dam.

Clays and carbonaceous minerals come from natural sources. The rates of C02 emission related to its obtaining are linked to the mining works of extraction and transport fundamentally. The processes of modification of the clays present low rates of C02

emission for being effected by electrical equipments.

Materials Oreanoclav. Also called, organophilic clay it is prepared with smectite clays, known as bentonite. The smectite clays are a group belonging to the group of phyllosilicates and its main feature is that there are those that swell in water or sodium clays and those that do not swell or swell little also called calcic clays. The smectites are represented by various types of clays such as montmorillonite. beydelite, nontronite and others. The montmorillonite is the principal member of the smectite group and its general formula is: (Na, Ca) 0.33 (AI, Mg) 2Si4O10 (OH) 2 • (H20) n. Were selected some calcic bentonite from the State of Paraiba, Brazil to modify them with sodium carbonate turning them into bentonite that swell in water or sodium clay. These clays modified by sodium were further modified using quaternary ammonium salt to produce organoclay which are characterized by being hydrophobic. This type of organoclay has shown physical and chemical compatibility with polymeric materials such as rubber. Used as reinforcing filler, the organophilic clay collaborate in structuring the nanocomposite at the stages of mass preparation and vulcanization, transferring to the final nanocomposite material technological properties equal or better than those presented by the materials traditionally prepared using carbon black as reinforcing filler.

Natural rubber. A natural product obtained from the Hevea Brasiliensis tree, it was provided by the Esper Ltd. The final mass formulation is referred in parts of fillers to 100 part of rubber or phr.

Chemical Additives for the process. This additives are used to provide better vulcanization process, less time, temperature and protection again climate, chemical and aging. These chemical are a) native sulfur (S) supplied by the company RCN 325 mesh, b) Zinc Oxide 99.9% purity supplied by the company BRASOXIDOS; c) stearin (stearic acid) supplied by Companhia

Paranaense de Sabonete & CHicerina. d) Accelerators for rubber TMTD (tetrametiltiurama disulfide) and MBT (2 Mercaptobenzotiazol) provided by FLEXYS.

Fillers. Carbon-black. During works for this paper it was used the type N660, or GPF provided by Cabot Co. such a choice is due to the fact that this type of carbon-black adds general-purpose resistance in the intermediate range for all devices where it is applied.

Additional reinforcing materials. Some different kind or carbonaceous minerals provided by Esper Ltd. [2]

Methods

Preparation of the mass composition.

The mass to produce the plates was prepared following the procedures used by the Esper Industries. For each assay was used 200 g of rubber as pattern reference. The fillers activators, accelerators and antioxidants were weighed in ratio according the Esper procedures. Before the vulcanizing process, the mixing process of substances was done in roller mixer brand MECANOPLAST model C250. The maximum temperature of the rolls was 45 ± 5 " C. The speed of the rolls was about 17 rpm and 15 rpm for the other. The speed difference of rotation between the cylinders allows breaking the structure of the rubber before carry out the mixture of the components of the elastomeric mass. After the mass of rubber being over ± 5 minutes under the action of the mixer cylinder, the structures of natural rubber have been broken, getting blanket form. Then it was started to add the reinforcing filler, then the activators, accelerators, and finally antioxidants. After approximately two minutes of action in the calender of the thus composed mass, the same one acquired homogeneous visual aspect (Fig, 1, above), being sent for vulcanization. There were prepared the following composition masses to be vulcanized: a) natural rubber (NR), b) natural rubber + carbon black 10 phr (CB10), c) natural rubber + carbon black 40 phr (CB40), d) natural rubber + organoclay 10 phr (OC10); e) natural rubber + carbonaceous discarded minerals 10 phr (CM10) and natural rubber + carbonaceous discarded minerals 10 phr (CM40).

Vulcanization of the mass.

The vulcanization was carried out using a form of steel under the action of a Thermo-Hydraulic Press from the Marconi Brand, model MA098, The mass was compressed to 15 tons of pressure to remove possible trapped gasses in the mass during the mixing process. The two plates of the press were heated to 160 ° C ± 5 ° C. The vulcanized rubber plates resulted after the vulcanizing time according the procedures established by the ESPER Industries (Fig. 1, below).

Conformation and characterization of the test samples.

For each composition two plates were shaped as a minimum, with dimensions of 150mm x 150mm x 2 mm. From each plate, three tie-shaped samples were cut for mechanical test. (Fig. 1).

The tie-shaped samples (Fig. 2) were made by cut on an OMAX -2652A System of Ultra-High Pressure Abrasive Water - Jet. This equipment is in the Mining & Petroleum Engineering Department at the Polytechnic School of the University of Sao Paulo, Brazil.

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Technological parameters, cost and time for cut was taken from succeeding procedures of the authors of this paper presented in international technical publications [6],

Those specimens were used to evaluate the mechanical properties of elongation of rupture and tensile strength according to ASTM D412 [7]. Hardness was evaluated in the rest of each plate according to ASTM 2240 - 05(2010) [8],

X - Rav Diffraction tests

Results

DRX- Results. Figures 3, 4, 5 and 6 present the DRX curves for main compositions of this work: Natural rubber, Natural rubber +

The XRD test was done to study the quality of interaction between rubber and reinforcing filler.

Mechanical tests: elongation of rupture and tensile strength and hardness

The elongation of rupture and tensile strength assays was developed in a Universal test machine model EM1C DL3000. Hardness, in a Sclerometer Hardness Tester.

carbon black (40 phr), Natural rubber + organoclay (10 phr) and Natural rubber + carbonaceous discarded minerals (40 phr)

Fig. 1 Above, masses after calendar process. Below, same masses after vulcanized.

Fig. 2 Tie-shaped samples cut by AWJ

385

Fig. 3 DRX curve of pure NR without fillers. Peaks to the right of the curve belong to chemical additives for vulcanization as S, TMTD, MBT and others.

Fig. 4 DRX curve of blanket made with NR + CB40 phr.

Fig. 5 DRX curve of blanket made with NR + OCIO phr.

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Fig. 6 DRX curve of blanket made with NR + CM40 phr.

Interpretation of the DRX results:

In the four curves it is possible to see the peak in 4,6 (or nearly) that correspond to the natural rubber demonstrating the integration or assimilation of the reinforcing filler by the rubber.

There is great similarity between DRX curves of Fig. 3 (NR) and Fig. 5 (NR-K)C10). There is some similarity between the curves of Fig. 3 (NR) and Fig. 6 (NR+CM40), but the less similarity

between curves is showed between Fig. 3 (NR) and Fig. 4 (NR+CB40).

Then, it might be seen that the resemblance between the curves can be an indicator of assimilation level or affinity between the rubber and the reinforcing filters. Thus, the more similar the curves of the plates of reinforced material with the curve pure of natural rubber, the greater the level of affinity or assimilation of the reinforcing filler by rubber. And the better the resulting reinforcement effect.

Sample composition NR NR + CB10 NR + CB40 NR + OC10 NR + CM10 NR + CM40 NR + OC10 + CM30

Table 1 Mechanical assay Rupture tension (MPa)

04,2 10,5 11,2 24,0 29,5 25,8 29,7

s effected to the samples Rupture lengthening (%)

675,0 415,0 315,0 625,0 638,0 757,0 767,0

Hardness Shore A 43 48 52 52

50,0 58,0 64

References

Conclusions

All the compositions reinforced with alternative fillers present superior performances in all the parameters than that reinforced with carbon-black. The composition of Natural rubber + lOphr organoclay + 30phr carbonous minerals show the betters results at all but not too distant from the others alternative fillers. For that reason it must be study in specials assays as abrasion resistance and permeability to define its performance in heavy requirement elastomeric vulcanized products.

1. Publication in Internet: Non-Black Fillers for Rubber The big 3. In http://www.rtvanderbilt.com/NonBlackFillers pdf. Accessed on October 2, 2011.

2. Fabio J. Esper. Substituiçào do Negro de Fumo por Composites e Nanocompôsitos de Borracha Natural. Tese de Doutorado. Escola Politécnica da Universidade de Sâo Paulo. 2010. 130 p.

3. Information on http7/education.;lab.or^'itseleinental/ele00b.hlml

4. G. R. Martin-Cortés, A. A. Silva, F. J. Esper, W. T. Hennies, F. R. Valenzuela-Diaz: NAOB® -Nanocompôsitos Argila OrganofilicafBorracha. In

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Proceedings of 19° CBECiMat - Congresso Brasileiro de Engenharia e Ciência dos Materials. CD-Rom. Campos de Jordäo. Brasil, Novembro - 2010.

5. Guillermo R. Martin-Cortes, Adriana A. Silva, Kleberson R. O. Pereira, Fabio J. Esper, Lisiane N. L. Santana, Wildor T, Hennies, Francisco R. Valenzuela-Diaz; Technology Characterization of Organo-Clays Obtained from Benfoniles of the State of Paraiha. Materials Science Forum Vols. 660-661 (2010) pp 1124-1129. © (2010) Trans Tech Publications, Switzerland.

6. W.T. Hennies, G.R. Martin Cortés, V.H. Lauand, CT. Lauand, A, Stellin Jr., G. De Tomi, F.R. Valenzuela-

Diaz. Applications of Ultra High-Pressure Powder-Abrasive Waterjet. In Materials Science Forum Vols. 660-661 (2010) pp 1137-1144. Online available since 2010/Oct/25 at www.scientific net © (2010) Trans Tech Publications, Switzerland. doi:10.4028/www.scienüfic.net/MSF.660-661.1137

7, ASTM D412. Standard Test Methods for Vulcanized Rubber and Thermoplastic Elastomers-Tension. Dec-2006.

8. ASTM D2240 Standard Test Method for Rubber Property - Durometer Hardness.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

PROPERTIES OF ADDITIONAL REINFORCERS MATERIALS USED TO COMPLEMENT NAOB® -A RUBBER / ORGANOCLAY NANOCOMPOSITE MATERIAL.

Fabio J. Esper; 'Guillermo R. Martin-Cortés; 3Luiz S. G. Dantas, 3Adriana Cutrim 'Wildor T. Hennies, Francisco R. Valenzuela-Diaz

"Mining & Petroleum Eng. Dept. - Polytechnic School - University of Säo Paulo Av. Prof Mello Moraes, 2373 Cid. Universitaria 05508-900 Säo Paulo, SP, Brazil.

2Raw Materials & Non-metallic Solids Laboratory - Metallurgie & Materials Eng. Dept. - Polytechnic School - University of Säo Paulo. Av. Prof. Mello Moraes, 2463 Cid. Universitaria 05508-900 Säo Paulo, SP, Brazil

3 BENTON1SA - Bentonita do Nordeste S. A. Joäo Pessoa Brazil

Keywords: discarded ultrafine carbonaceous minerals (1), vulcanized rubber products (2), reinforcing raw materials for rubber products (3), organic modified clays (4), nanocomposite material with organophillic clay and additional reinforcers materials (5)

Abstract

Vulcanized rubber products are widely spread. Auto parts, tires, sport shoes and other products are examples of it. Most of the raw materials used in the formulations of vulcanized rubber products come from industrial process consuming time and energy but also generating C0 2 emissions during the respective transforming processes. The traditional main reinforcing material used in rubber formulations is carbon-black which is a petroleum derivative obtained by the incomplete burning of the fossil fuel. The NAOB - a rubber / organoclay nanocomposite material was developed to make vulcanized rubber products without carbon-black. In NAOB the primary reinforce filler is the organoclay, which participates in the formulation of the nanocomposite in much smaller quantities than the carbon-black. So to keep the material specific volume, maintaining or increasing its technological characteristics, additional reinforcing materials are used. This article presents some of the additional reinforcing materials for the NAOB and their technological characteristics.

Introduction

Commonly, rubbers vulcanized products show black or a very dark color. Basically, the formulations of vulcanized rubber, whether natural or synthetic rubber, are evaluated on the following laboratory tests: elongation; tensile properties; tensile strength; tensile stress, abrasiveness resistance, resilience, and impermeability among others depend in high ratio of the type of reinforcing filler added, traditionally, to the several carbon-black classes and their mixing proportions.

There are other reinforcing materials for rubber whose colors are not dark but clear other than black. In a general way, the most common non-black fillers for rubber are calcium carbonate, kaolin clay, precipitated silica, talc, barite, wollastonite, mica, precipitated silicates, fumed silica and diatomite. Of these, the three most widely used, by volume and by functionality, are calcium carbonate, kaolin clay and precipitated silica [1 ].

By the way, in the few last years, researchers from two departments of the Polytechnic School of the University of Sào Paulo find that there are other type of reinforcing filler materials instead carbon-black and those non-black cited above. When these new reinforcing fillers materials are used the vulcanized rubber products present different appearance but also its technological performances are enhanced. As these materials are used in fewer quantities then carbon-black researchers seek complement materials to obtain same volumes of production. In this research

line it was find that some nanoparticles materials from modified non-metallic minerals and also others nanoparticles substances from different sources, alone or mixed among them, used in substitution of carbon-black as rubber reinforce filler are showing, an enhancing technological behavior in comparison with the traditional carbon-black or a similar behavior [2].

The characteristics which determine the properties that filler will impart to a rubber compound are: particle size, particle surface area, particle surface activity and particle shape. Surface activity relates to the compatibility of the filler with a specific elastomer and the ability of the elastomer to adhere to the filler. Functional fillers transfer applied stress from the rubber matrix to the strong and stiff mineral. It seems reasonable then that this stress transfer will be better effected if the mineral particles are smaller, because greater surface is thereby exposed for a given mineral concentration. And if these particles are needle-like, fibrous or platy in shape, they will better intercept the stress propagation through the matrix. A compound's physical/mechanical properties can be strongly influenced by the surface activity of the filler, which is the ability of the filler's surface to bond with the matrix

m Carbon-black, material derived from oil is produced during the refining of crude oil or mineral oil with the known consequences of environmental contamination. Its main physical properties are nanometric grain sized distribution and a depth black color. It is an amorphous carbon formed when the hydrocarbon is burned without enough oxygen for it to burn completely. It's also known as lampblack, gas black, or channel black, is used to make inks, paints and rubber products. It can also be pressed into shapes and is used to form the cores of most dry cell batteries, among other things [3]. The substitution of this substance, traditionally used as filler to reinforce the vulcanized rubber products, will help to reduce dependence on petroleum products and also to reduce environmental contamination.

NAOB® a rubber + organoclay nanocomposite [4] is a patented new type of material where smectite clay modified by ammonium quaternary [5] through low wet procedures and some additional materials, replaces carbon black as filler. Smectite clay has being used in several purposes, due to its excellent physical, chemistries, and structural properties. Main applications of smectites in Brazil are as agglomerative in métallurgie molds, in iron ore pelleting, as drilling mud in oil and water catch wells, as clarifier of wines, in catalyses at the chemistry and pharmaceutical industries, for pollutants removal in water treatment plants, in cosmetic among other applications.

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In the composition of vulcanized rubber products is used, at least, 40 phr of carbon-black. It is not uncommon rubber compositions prepared for vulcanized with 50, 60 or more phr of carbon-black. As reinforce fillers NAOB ® uses only 10 phr organoclay + 30 phr of additional materials, to achieve similar volumes of products to those obtained with carbon-black. Those additional materials could come from several different sources: from minerals with small or without industrial application, chemically modified or not, even different types of materials discarded by industry, also including those minerals discarded in the process of mineral extraction or end-discarded products as ashes from energy production [4].

This paper describes some of the additional reinforcers materials from several sources, natural or anthropogenic, used in NAOB® to produce the same volumes of material obtained in vulcanized rubber through traditional ways of production with carbon-black as filler. Also describe its properties evaluated by laboratory tests on XRD - X ray diffraction and SEM - Scanning Electronic Microscopy. Actually, some of these additional materials are minerals or mining by-products with poor or without economical applications. Some of those materials have not industrial applications because its physical properties are not adequate for the industries at the region around. Today destinations of those materials are: abandoned in their own field without extracting or the mine tailings dam or the landfill with obvious bad consequences to the environment. Finally, this paper complement another one present in this same event and that describe the behavior of NAOB® in some mechanical laboratory test.

Clays and carbonaceous minerals come from natural sources. The rates of C02 emission related to its obtaining are linked to the mining works - extraction and transport - fundamentally. Processes for clays modification have low rates of C02 emission because it is developed by electrical equipments. Anthropogenic materials referred to are those discarded by the industry in landfills whose C0 2 emissions are related to transport to the recycling facilities.

Additional reinforcing materials, a) Argillaceous abandoned materials (Aam) or discarded at the mine tailing dam from the smectite clay mining activities slightly modified with salts of ammonium quaternary; b) some different kind of carbonaceous minerals (Cms) and c) those discarded as waste by the industries as Rice Husk Ashes (RHA) and Cellulose (Cel) obtained from the disposal of parts of trees, provided by Esper Ltd. [2]

Methods

X - Ray Diffraction It was used the Philips diffractometer model XPert MPD equipped with copper electrode radiation. The technical data of equipment are: Tube anode: Cu; Voltage [KV]: 40; Current [mA]: 40; Wavelength: 1.54060; Time per step [s]: 1.000. The tests were performed to samples in powder form.

IR - Infrared spectroscopv Analyses of infrared spectroscopy were performed in Magma FTIR spectrometer brand ESPS Nicolet model 560, coupled to a microcomputer with records of the spectral range 4000-400 cm-l, with a resolution of 4 cm-l, making 64 sweeps, purging with compressed air, with enhancement of transmission. The powder samples were prepared and placed in the sample holder and analyzed by the method of diffuse reflectance.

SEM - Scanning Electronic Microscopy It was used scanning electron microscopy PHILLIPS brand, model XL30 EDAX. The samples were attached on substrates of aluminum and coated with gold or graphite in a vacuum chamber, using sputter Balzers Model SCD 050 Brand BAL-TEC.

Results

The following images (Fig. 1, 2, 3, 4) show results obtained in the X - Ray Difraction of the materials studied in this paper.

Fig. 1 XRD curve from the argillaceous abandoned materials from the smectite clay mining activities with slightly modified with salts of ammonium quaternary. [6]

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Fig. 2 XRD curve of the rice husk ashes with it typical Si peak.

Fig. 3 XRD curve of a carbonaceous mineral.

Fig. 4 XRD curves of the cellulose from the disposal of parts of trees.

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Interpretation of the XRD results same area than the smectite clay. The peaks of Si (quartz) it is practically present in all XRD curves.

The Fig. 1 shows that the modification with the ammonium salt managed a little change at the inter-planar distance of the basic SEM - Scanning electronic microscope micrograph. cell of the argillaceous material. Fig. 2 presents the typical quartz peak at 23 angle. Fig, 3 shows some peaks possible from some The following images (Fig. 5, 6, 7) present the micrograph and the argillaceous impurities. Fig. 4 Show a great peak at an angle in the X - ray microanalysis of the materials studied in this paper.

Fig. 5 SEM micrograph: a) argillaceous abandoned materials from the smectite clay mining activities after slightly modification with salts of ammonium quaternary; b) carbonaceous minerals; c) Rice Husk Ashes and d) Cellulose from the disposal of parts of trees.

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Fig. 6 Rice husk ashes: Si is the main element, C, 0 and K have smaller peaks.

Fig. 7 Cellulose from the disposal of parts of trees: Again Si is the ma with less size peaks are also present

Interpretation of the SEM Micrographs a and b were take with 35000 and 50000 increases. Both micrographs present nanoparticles in thin layer shape. Micrograph c of RHA show a mix of fibers, particles and nanoparticles, meanwhile micrograph d present "large" fibers with a few nanosized thin layer particles.

element. Cl, C and AI present medium size peaks and Na, Ca e Fe the composition of this material.

In both SEM - X-ray Microanalysis graphics, the Si is the main element, but the presence of other elements as C, AI, Cl, Na, Ca e Fe could explain why these materials collaborate with the organoclay to transmit better mechanical properties to the rubber vulcanized material.

Conclusions

The shape, size and structure of the particles and fiber components of this studied additional materials and its chemical composition explain why the combination of it with the organoclay contributes to improve the mechanical properties of vulcanized rubber products. All this explains why the NAOB surpasses technologically in many respects to the products made with carbon black as reinforcing filler.

References

1. Publication in Internet: Non-Black Fillers for Rubber The big 3. In http://www,rtv.anderbilt com/NonBlackFillers pdf. Accessed on October 2, 2011.

2. Fabio J. Esper. Substituiçào do Negro de Fumo por Composites e Nanocompositos de Borracha Natural. Tese de Doutorado. Escola Politécnica da Universidade de Sâo Paulo. 2010. 130 p.

3. Information on hftp://education.)lab.or^/itselemental/ele006 html

4. G. R Martin-Cortes, A. A, Silva, F. J. Esper, W. T. Hennies, F. R. Valenzuela-Diaz: NAOB® -

Nanocompositos Argila Organofilica/Borracha. In Proceedings of 19° CBECiMat - Congresso Brasileiro de Engenharia e Ciência dos Materiais. CD-Rom. Campos de Jordâo. Brasil. Novembro - 2010.

5. Guillermo R. Martin-Cortes, Adriana A. Silva, Kleberson R. O. Pereira, Fabio J. Esper, Lisiane N. L. Santana, Wildor T. Hennies, Francisco R Valenzuela-Diaz: Technology Characterization of Organo-Clays Obtained from Bentonites of the State of Paraiba. Materials Science Forum Vols. 660-661 (2010) pp 1124-1129. © (2010) Trans Tech Publications, Switzerland.

6. Guillermo R. Martin-Cortes, Fabio J. Esper, Francisco R Valenzuela-Diaz. Relatörio Aprovado correspondente à Fase I do Projeto PIPE-NAOB -Projeto de Inovaçâo Tecnotôgica em Pequenas Empresas Brasileiras - Nanocompositos de Arguas Organofilicas / Borrachas, apresentado à FAPESP -Fundaçâo de Amparo à Pesquisa do Estado de Säo Paulo. 2008. [Report Approved corresponding to Phase I of the Project NAOB-PIPE - Project of Technological Innovation in Brazilian Small Business - organoclay nanocomposites / Rubber presented to FAPESP -Foundation for Research Support of Säo Paulo 2008]

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THERMAL PROPERTIES OF CARBON NANO TUBES REINFORCED MG-MATRIX NANOCOMPOSITES

Sardar S. Iqball2, Aydin Mustafa24, Saikat Talapatra3, Peter Filipl2

'Department of Mechanical Engineering & Energy Processes, S1U Carbondale, IL, 62901 2Center of Advanced Friction Materials, SIU Carbondale, IL, 62901 4University of Dumlupinar, Turkey

'Department of Physics, SIU Carbondale, IL, 62901

Abstract

The present work is the study of thermal properties of pure Magnesium (Mg) and Mg-CNT reinforced with virgin carbon nano tubes (CNT) in the temperature of 25-500°C. The Mg and Mg-CNT nanocomposites reinforced with CNTs at the ratio of 0.5, and 1 % by weight were prepared with powder metallurgy method by cold pressing followed by sintering at 620 °C for 2 hours in argon atmosphere. The microstructure was characterized by polarized light microscopy and SEM, and structure by TEM and X-ray diffraction patterns. Density of pure Mg decreased with increasing CNTs whereas porosity increased. Thermal diffusivity and heat capacity of 1%CNT-Mg nanocomposites show highest values than pure Mg. Similarly thermal conductivity of 1%CNT-Mg nanocomposites was the highest. The results show that thermal properties of Mg improved significantly with the addition of CNTs.

Keywords: Magnesium, CNTs, powder metallurgy, thermal properties, SEM, TEM.

Introduction

In recent years, studies have been carried out on the development of carbon nanotubes (CNT) reinforced metal matrix composites due to CNT more superior mechanical and thermopysical properties to other reinforcement materials. Magnesium based metal matrix composites (MMC) have a great interest due to their low density (1.738g/cm3) and high specific properties [1]. However pure Magnesium is not suitable for service at 150-300oC resulting in lower mechanical strength. The values of the CTE and thermal conductivity of Mg at room temperature are 25xl0-6K"' and 156 Wm'lC' [2]. The coefficient of thermal expansion of alloys and metal matrix composites (MMCs) is an important thermomechanical property. Structural components in service are subjected to temperature changes and it is important to minimize dimensional changes. Thermal conductivity plays an important role in the performance of alloys and MMCs in structures. The value of thermal conductivity is an important criterion for the selection of alloys for various applications. Numerous studies are available about mechanical properties of nanocomposites [3-6], however, very few studies provide the information about thermal properties of CNTs reinforced Mg-based nanocomposites.

Powder metallurgy method has been the most successful process for infiltration of metal into the preform of nanotubes [7]. Several studies reported [7-9] about the successful

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incorporation and to homogenize the mix of CNTs in aluminum (Al) and magnesium (Mg) matrices by ball milling in Turbula mixer by powder metallurgy techniques. The aim of the present study is to determine thermal conductivity by measuring coefficient of thermal expansion (CTE), thermal diffusivity (a) and specific heat capacity (Cp) of Mg nano composites in a temperature range of 25-500°C.

Experimental Procedures

Pure magnesium powder (99.9% pure, -325 mesh in size) from Alfa Aesar was used as metal matrix. The multi wall carbon nano tubes (MWCNTs) produced by chemical vapor deposition (CVD) process were used as reinforcements for the fabrication of nanocomposites. Mg-nanocomposites were prepared by mixing CNTs at ratio of 0, 0.5, and 1% by weight. The mixing was performed by high energy ball milling machine (Turbola spex 8000) for 5 minutes. The mixture of Mg and CNTs was cold pressed at 700 MPa. Mg-nanocomposites were sintered at 620 °C under argon atmosphere for 2 hours in tube furnace. The nanocomposites pure Mg, Mg-0.5wt%, Mg-lwt% were named as Mg, 0.5Mg, and lMg based on their weight percentage in Mg matrix, respectively. The density was measured by mineral spirit method using Archimedes' principle before and after sintering. Thermal diffusivity was measured by using laser flash method, NETZSCH DIL-402C was used for measuring coefficient of thermal expansion and heat capacity was obtained by NETZSCH DSC-404C. Microstructure and morphology of the polished and fractured surfaces were analyzed using Nikon Microphot polarized light microscope and scanning electron microscopy (SEM, Hitachi S570), respectively. The structures were characterized by Hitachi H7650 transmission electron microscopy (TEM).

Results and Discussion

Microstructure

Figure 1 shows the TEM images of multi walled carbon nano tubes (MWCNT), which are used in the current study. The diameter of CNT is approximately 100 nm. It is apparent that CNTs consist of amorphous and crystalline structures. CNTs appear to have catalytic particles in them.

Fig.l. TEM images of CNT. The micron bar is 500 nm.

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The presence of CNTs in the microstructures was not detected by XRD analysis and it could be due to the limitation of filtered X-ray radiation to detect the phases with less than two volume percent in a multiphase structure as was also reported by Aung et al. [10].

Figure 2 shows the microstructure of Mg-CNT nanocomposites for 0.5 and 1% CNTs. The distribution of CNT is very low and appears to have agglomerated for 0.5Mg nanocomposite. On the other hand, lMg nanocomposite shows the appreciable dispersion of CNT within the Mg-matrix despite occasional clustering of CNTs.

Fig. 2 Microstructure of Mg-CNT nanocompositesf for 0.5 (a), and 1% CNT (b), respectively. Micron-bars is 8 um.

Density and Porosity

Figure 3 shows calculated and measured density, and Fig. 4 shows porosity of nanocomposites after sintering. The density of nanocomposites steadily decreases with pure Mg exhibiting the highest density and lMg-nanocomposites the lowest. Similarly, the porosity progressively increases with pure Mg-exhibiting the lowest porosity and IMg-nanocomposite the highest porosity. The steady decrease in density and increase in porosity is due to the increase in the ratio of CNTs. This indicates that although, the decrease in density has occurred due to CNT reinforcements, yet an appreciable increase in porosity is observed. The density of 0.5Mg-nanocomposite is not so drastic when compared to the pure Mg, which indicates a negligible effect of the addition of CNTs at 0.5 %wt in Mg matrix, however, the porosity doubles after the addition of CNT into Mg matrix. However, with the addition of 1%CNT, the decrease in the density is 1.4% of pure Mg, whereas the porosity increases 28 times than that of pure Mg. This indicates the less effective homogeneous distribution of the CNTs in Mg-matrix, as the clustering of the CNTs is unavoidable due to the sheer size of the CNTs at the nano levels, since the mixing time was short (5 minutes) in order to avoid crushing of the nanotubes into a much smaller sizes .

However, Aydin et al. [11], Gupta et al. [12], and Wang et al. [13] reported an increase in density of Mg-based naocomposites after the addition of SiC and TiC nanoparticles in comparison to pure Mg. For lightweight applications of Mg-based composites, it is highly desirable to have low density Mg-nanocomposites which is evident in the case of present study. It is pertinent to note that although measured density is in accordance with theoretical values, yet the corresponding measured values of density of nanocomposites are less than their

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corresponding theoretical values. The density was measured for nanocomposites before and after the sintering in order to observe the effect of the addition of CNTs in Figure 3. It shows that before sintering the density of nanocomposites consistently decreases with the increasing amount of the CNTs.

Fig.3 Density(a), and porosity(b) before and after sintering of Mg-nanocomposites.

However, the density after sintering decreases for nano composites with higher CNT. On the other hand, a reciprocal effect is observed in terms of porosity where it increases with the increasing CNTs in Mg-matrix, which is in agreement with the reported literature [14, 15]. This could be mostly due to the clustering effect of the CNTs that induce the empty spaces and voids resulting in higher porosity. It is pertinent to note that the porosity increases after sintering when compared to the porosity before sintering. It could be attributed to the expansion of nanopores in nanocomposites before sintering into much bigger micro and macro pores after sintering and resulting in creating a network of micro to macro pores. The sintering at higher temperature causes the Mg-matrix to expand more than CNTs, the mismatch between thermal expansion coefficients puts the interface under a very high stress between matrix and CNTs which is compressive in nature. On the other hand, a vice versa effect is observed on cooling of the nanocomposites. The Mg-matrix shrinks more than CNTs and thus puts a very high stress at the interface between matrix and the CNTs which tends to be tensile in nature. The mismatch between thermal expansion coefficients also results in creating a network of porosity ranging from micro to macro in nature. The nonopores can be converted into micropores by bridging of the submicron pores after sintering.

Thermal Properties

Coefficient of Thermal Expansion

Figure 4 shows the coefficient of thermal expansions (CTE) of Mg and Mg-CNT. CTE of all the specimens is decreasing with temperature; with CTE of pure Mg is the highest. CTE of nanocomposites show a marked decrease with increasing temperature with Mg-0.5%CNT nanocompsites exhibiting lowest thermal expansion closely followed by Mg-1%CNT nanocomposites, which could be attributed to the low amount of CNT (0. 5 wt%) and possibility of good dispersion in Mg-matrix with less of agglomeration, due to which the well dispersed CNT could have been able to constrain the natural expansion of Mg-matrix due to higher surface area of CNT. On the other hand, CTE of 1%CNT-Mg nanocomposites is higher than the 0.5% CNT-Mg composites but lower than pure Mg, which could be due to the more of agglomeration

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Fig. 4 Thermal expansion of Mg and Mg-CNT nanocomposites with temperature.

of CNTs that provides less surface area due to lower dispersion of CNTs in the Mg-matrix as a result the Mg-matrix expansion is slightly hindered.

Specific Heat Capacity

Figure 5 shows the heat capacity curves for pure Mg and Mg-CNT nanocomposites with respect to temperature up to 525°C. It is apparent that the heat capacity of pure Mg is lower than the nanocomposites during the test heating run. With the addition of 0.5 and 1 wt% CNT, the heat capacity of nanocomposites improves significantly. Heat capacity of lwt% CNT nanocomposites is the highest and continues to increase with temperature with slight dip in its capacity around 420°C, but then shows an increase afterward. The heat capacity of 05. wt% CNT nanocomposites is higher than pure Mg, but lower than 1 wt% CNT nanocomposite and exhibits identical trend as 1 wt% CNT nanocomposites. It indicates a significant beneficial effect of adding a very small amount of CNT on the heat capacity of Mg-matrix composites.

Fig. 5 Heat capacity of Mg and Mg-CNT nanocomposites with temperature.

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Thermal Diffusivity

Figure 6 shows the plots of thermal diffusivity of pure Mg and nanocomposites with respect to temperature.

Fig. 6 Thermal diffusivity of Mg and Mg-CNT nanocomposites with temperature.

Thermal diffusivity of pure Mg is about 0.9 cm2/s at room temperature. As the temperature increase thermal diffusivity of pure Mg starts to decrease and shows a linear trend after 200 oC till maximum temperature of 500 oC. On the other hand, thermal diffusivity has improved for nanocomposites with the addition of 0.5 and 1 wt% CNT in Mg-matrix. However, thermal diffusivity of nanocomposites decreases with test heating run but is still higher than pure Mg during the test heating run, with 1 wt% CNT nanocomposites showing higher thermal diffusivity than 0.5 wt% CNT nanocomosites. However, rate of decrease of thermal diffiisivity of 1 wt% CNT nanocomposites smaller than 0.5wt% CNT nanocomposites, and indicates a relative improvement in thermal property of Mg-matrix nanocoposites.

Thermal Conductivity

Thermal conductivity is derived by measuring the density, heat capacity and thermal diffusivity of the nanocomposites give by following equation;

k = c p * a * p (1)

Figure 7 shows thermal conductivity of pure Mg and Mg-matrix nanocomposites with temperature. It is obvious that thermal conductivity of pure Mg has improved significantly with the slight addition of CNTs. At room temperature, thermal conductivity of pure Mg is lower than the nanocomposites. Thermal conductivity of pure Mg increases significantly with the addition of 0.5 and 1 wt % CNT in Mg-matrix. Thermal conductivity of pure Mg and Mg-nanocoposites decreases with temperature. However, there is a marked decrease in thermal conductivity of 1 wt% CNT Mg-nancomposites around 300oC. An overlap of thermal conductivity of 0.5 and 1 wt% CNT Mg-nanocomposites appear to occur with thermal conductivity of lwt% CNT Mg-nanocomposites decreasing than 0.5 wt % CNT Mg Nano composites. This trend continues till maximum temperature. An interesting observation is noted in case of thermal expansion of nanocomposites. It is obvious that 0.5 wt% CNT nanocomposites exhibit the lowest coefficient of thermal expansion than pure Mg and 1 wt% CNT nanocomposites. Since the density of 1%

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CNT nanocomposite will be decreasing due to expansion of volume as compared to 0.5 wt% CNT nanocompositres. This could have impacted thermal conductivity of 1 wt% CNT nanocomposites since large thermal diffiisivity and conductivity are attributed to the higher density of the sample [16]. However, thermal conductivity of pure Mg improved 25 and 55% with the addition of 0.5 and 1 wt% CNTs, respectively, in Mg-matrix at room temperature. Similarly, thermal conductivity of 0.5 and 1 wt % CNT was 49 and 38% higher than pure Mg at maximum temperature during test heating run, indicating that by employing a small amount of CNT, thermal conductivity of pure Mg could be enhanced significantly.

Fig. 7 Thermal conductivity of pure Mg and Mg-CNT nanocomposites with temperature.

Conclusions

1. The fabrication of Mg-MWCNTs nanocomposites was successfully accomplished by high energy ball milling powder metallurgy method.

2. The density of nanocomposites decreases before and after sintering with the increasing amount of the CNTs. The increase in porosity before and after sintering could be attributed to the expansion of nanopores into a network of much bigger micro and macro pores.

3. CTE of pure Mg and nanocompsites decreases with temperature. Heat capacity of pure Mg is lower than the nanocomposites, which shows a beneficial effect of adding a very small amount of CNT on the heat capacity of Mg-matrix composites.

4. Thermal diffiisivity of pure Mg is the lowest. An increase in thermal diffiisivity of nanocomposites with the addition of 0.5 and 1 wt% CNT in Mg-matrix is achieved.

5. Thermal conductivity of pure Mg improved 25 and 55% with the addition of 0.5 and 1 wt% CNTs, respectively, in Mg-matrix at room temperature. Similarly, thermal conductivity at maximum temperature was 49 and 38% higher for corresponding nanocomposites than pure Mg, indicating that by employing a small amount of CNT, thermal conductivity of pure Mg can be enhanced significantly.

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References

[I] ME. Alam, S. Han, Q. B. Nguyen, A. M. Salem, M. Gupta, "Development of new magnesium based allys and their nanocompsoites," J. of Alloys & Compounds, 509 (2011), 8522-29.

[2] Z.J.Min, Y.Ying, M.Lamvik, W.Gang, "Determination of thermal conductivity of Magnesium-alloys," ./.Onr. South Univ. Techno/., 8 (2001), 60-64.

[3] A. M.K. Esawi, M. A. El Borady, "CNT Reinforced Aluminum Strips," Comps. Sei. & Tech 68 (2008), 486-92.

[4] C.S. Coh, J. Wei, L.C. Lee, M. Gupta, "Development of novel CNT reinforced magniziuin annaocomposites uising the powder metallurgy technique," Nanotechnology, 17 (2006), 7-12.

[5] A. Esawi, K.Morsi, "Dispersion of CNTs in aluminum powder," Composites: Part A 38 (2007), 646-60.

[6] Y. Feng, H. Long Yuan, Min Zhang, "Frabrication and properties of Ag-matrix composites reinforced by CNTs,"M7(eràfe Characterization, 55 (2005), 211 - 218.

[7] A. Esawi, K.Morsi, "Dispersion of carbon nanotubes (CNT) in aluminium powder," Composites: Part A,i% (2007), 646-60.

[8] [ 9 ] Carreno-Morelli E, Yang J, Couteau E, Hernadi K, Seo JW, Bonjour C, et al. Carbon nanotube/magnesium composites. Phys Stat Sol 2004; 201(8):R53-5.

[9] S. S. Iqbal, M. Aydin, S. Talapatra, P. Filip, R. Koc, "Effect o f Processing Duration on Mechanical Properties of Mg-based Nanocomposites Reinforced by Multi Walled Carbon Nano Tubes (MWCNT)," In. Proceeding of 6th International Conference on Powder Metallurgy & Exhibitions 2011, Oct 5-9,2011, Ankara, Turkey.

[10] N.N. Aung, W. Zhou, C.S. Goh, S.M. Ling Nai, J. Wei, "Effect of CNT on corrosion of Mg-CNT composites," Corrosion Sei. 52 (2010), 1551-53.

[II] M. Aydin, Rasit Koc, "Properties of Nano-Metal Carbide Contained Mg-TiC (SiC) 147," In Proceedings: 34th ICACC, Nanostructured Materials & Nanotechnology IV, January 24-29,2010, Daytona Beach, FL.

[12] M.Gupta, "Syntheis, microstrucutre, and properties characterization of disintegrated melt deposited Mg-SiC Composites," J. of Mat. Sei. 35(2000), 2155-65.

[13] H.Y. Wang, Q. C. Jiang, X..L. Li, J.G.Wang, Q.F.Guan, H.Q.Liang, "In situ synthesis of TiC from nanopowders in a molten magnesium alloy," Materials Res. Bulletin, 38 (2003), 1387-92.

[14] C.S. Coh, J. Wei, L.C. Lee, M. Gupta, "Development of novel CNT reinforced magnesium nanocomposites using the powder metallurgy technique," Nanotechnology, 17 (2006),7-12.

[15] R. P. Bustamante, I. E. Guel, W. A. Flores, M. M. Yoshida, P.J. Ferreira, R. M. Sanchez, "Novel Al-matrix nanocomposites reinforced with MWCNTs," J. of Alloys & Compounds, 450 (2008), 323-326.

16] L. Kumari, T. Zhang, G.H. Du, W.Z. Li, Q.W. Wang, A. Datye, K.H. Wu, 'Thermal properties of CNT-alumina nanocomposites," Composites Sei. & Tech., 68 (2008), 2178-83.

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TIMIS2012 141 s t Annual Meeting & Exhibition

New Advances in Synthesis, Characterization, and

Application of Layered Double Hydroxides

Edited by:

Jewel Gomes

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DESIGNING LAYERED DOUBLE HYDROXIDES FOR TARGETED APPLICATIONS

Jeanne M. Hossenlopp. Stephen Majoni, Cleopas Machingauta

Marquette University Department of Chemistry PO Box 1881, Milwaukee, WI 53201-1881, USA

Keywords: layered double hydroxides, hydroxy double salts, structure-reactivity relationships

Abstract

Layered double hydroxides (LDHs), and other similar nanodimensional layered metal hydroxides, offer promise for a number of applications ranging from controlled release of bio-active compounds to polymer fire retardant additives. Tailoring LDHs for controlled release or adsorbant applications requires an understanding of the kinetics and thermodynamics of anion release or analyte uptake. Recent work in our laboratory on the effects of systematic alteration of the metal hydroxide layer and/or anion structure will be discussed. The composition of the metal hydroxide layer provides wide tunability in reactivity for controlled release or anion exchange reactions. Anion polarity and hydrogen bonding capability can be utilized to vary the extent and rates of exchange reactions. Similar approaches to LDH structural alterations are applied to the design of potential polymer fire retardant additives. The effect of metal hydroxide layer composition and anion structure and packing influences the morphology of LDH-polymer (nanocomposites) and their thermal/fire properties.

Introduction

Layered metal hydroxides with exchangeable interlayer anions (An) can be utilized to provide nanodimensional structures with tunable physical and chemical properties. Examples of these materials include layered double hydroxides (LDHs) [M2+

XM3\.X (OH)2]!'+(Anx/11)xmH20,

hydroxy double salts (HDSs) [(M2+i.xMe2Vx)(OH)3(i.yyn]An-(1+3yynxmH20 where M2+ and Me2+

represent the different divalent metals, and the analogous layered hydroxy salts (LHSs) containing a single metal. The ability to vary the intralayer metal composition as well as the interlayer anion identity makes these materials attractive for a broad range of applications including use as catalysts [1,2], catalyst precursors [3,4,5], catalyst support materials [6], adsorbents, antacids, and ion exchangers [7,8]. A key difference between LDHs and layered metal hydroxides with exclusively divalent metals (i.e. HDSs and LHSs with a divalent metal) can be the nature of the anion interaction with the metal hydroxide layers [9].

Optimizing these materials for targeted applications requires the ability to control physical and chemical properties by utilizing variations in the metal hydroxide layer as well as the inter-layer anion content and packing. Different synthetic routes for obtaining, and modifying, these materials have been reported. A common strategy is to first synthesize a structure with a simple, easily exchanged anion, such as nitrate, and then to exchange the nitrate with another species

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with desired size, polarity, and/or other variation in physic-chemical properties. For example, there have been several studies that characterize anion-exchanged HDS materials [10]. Similarly, anion exchange reactions have been utilized with single-metal-containing layered hydroxyl salts, LHSs, such as intercalated copper or nickel hydroxides, M2(OH)3X, where X is an intercalated anion [11].

While layered metal hydroxides have been shown to be useful in a variety of applications, there are still many questions about the factors that control the basic reactivity of anion exchange reactions, a primary tool in optimizing structures. Work in our laboratory focuses upon exploring the factors that control structure-reactivity properties in model LDH and HDS systems.

Experimental

A precursor HDS, zinc copper hydroxy acetate (Zn/Cu-Ac), was prepared according to a literature method [10(b)]. 0.41g of ZnO in 5 ml of deionized (DI) water was added slowly to a solution of 1.00 g of Cu(CH3COO)2-H20 in 5 ml of DI water over a period of about 10 minutes. The resultant mixture was stirred frequently at room temperature for 24 hours. ZC-o-HCn was prepared by mixing 20.0g Of ZC-Ac with 1000 cm3 of a 0.1 M o-hydroxycinnamate (o-HCn) solution at room temperature for 24 hours with frequent stirring; the exchange reaction was carried out twice using the same conditions. ZC-w-HCn was prepared by reacting ZC-Ac with 0.5 M m-hydroxycinnamate (/w-HCn) solution at 40°C for 24 hours with frequent stirring, the exchange was carried out three times under the same condition each time. 1.0 Mp-HCn was used in the preparation of ZC-p-HCn with the exchange reaction being done once at 40°C, the higher concentration was necessary to obtain a single p-HCn phase.

Hydroxy nitrates were prepared via a similar route. A volume of 60 mL of 1 M Ni(N03)2 solution (0.06 mol) was added to 9.8 g of ZnO (0.12 mol) dispersed in about 10 mL of de-ionized water, and the mixture was stirred vigorously for about 10 mins. The mixture was then allowed to stand for 6 days at 65 °C. The solid was isolated, washed using warm deionized water, and dried at room temperature for 48 hrs. Acetic acid solutions for exchange reactions were prepared by mixing sodium hydroxide and acetic acid in a 1:1 mole ratio. The pH of the solutions was adjusted to 5.6 ± 0.2 using dilute NaOH and HCl. Finally, other acetate-containing HDSs were prepared as noted about for Zn/Cu-Ac.

The solid samples were analyzed using on a Rigaku Miniflex II diffractometer using Cu Ka (X =1.54 Â) as the radiation source operated at 30 kV and 15 mA. The powdered samples were pressed onto a glass sample holder until the surface was smooth. The diffractograms were recorded from 2.5° to 45.0° 20. Infrared spectra were obtained on a Perkin Elmer Spectrum 100 FT-IR spectrometer. The IR spectra were recorded using a MIRacle, Single Reflection Horizontal ATR Accessory with a ZnSe prism, the spectral range was 4000 cm"1 - 675 cm"1, and a scan resolution of 2 cm'1.

Results and Discussion

General Reactivity Trends: Hydroxy Cinnamate Release

The trends in anion exchange rates depend on the details of the combination of anions involved in the exchange as well as the composition of the metal hydroxide layer. The release of three

406

isomers of hydroxy cinnamatic acid via anion exchange provides an initial overview of the complexity of predicting reactivity trends.

</° HC.

% CH

OH

^ ^

(a) (b)

Figure 1. Structures of the anions used in this study: (a) o-hydroxycinnamate, (b) m-hydroxycinnamate, (c)/J-hydroxycinnamate. They are all in the trans-isomer as supplied by the

supplier, Sigma Aldrich.

For an initial comparison, a series of layered metal hydroxides with copper (CuH), zinc and copper (Zn/CuH), zinc (ZnH) and zinc and nickel (Zn/NiH) were synthesized with each of the three hydroxycinnamate isomers. Initial exchange reactions with chloride anions were carried out at 40 °C. Table I provides a qualitative assessment of the reactivity trends, determined via inspection of powder x-ray diffraction analysis of solid products remaining after 24 hours.

System Cu-o-HC Cu-m-HC Cu-p=HC Zn/Cu-o-HC Zn/Cu-m-HC Zn/Cu-p-HC Zn-m-HC Zn/Ni-m-HC Zn/Ni-o-HC

Table I: General Reactivity Exchange ion Chloride Chloride Chloride Chloride Chloride Chloride Chloride Chloride Chloride

Trends Exchanged? Yes yes (partial) No Yes yes (partial) yes (partial) No No No

In general, the para isomer is difficult to remove from the layers with the halogen or formate exchange partners. However, preliminary studies with longer alkyl carboxylate exchange partners indicate that as much as 87% of the CuH-p-HC precursor can be exchanged with octanoate under conditions where little-to-no reaction is observed using chloride as the exchange partner. The Zn/Cu HDSs exhibited the most consistent pattern of release of the hydroxy cinnamate isomer by exchange with chloride; similar results are obtained using bromide or formate anions.

407

Detailed Investigation of Hvdroxv Cinnamate Release from Zn/Cu HDS

The survey work presented above suggests that the structure of the anion is important in determining whether or not it can be exchanged under a given set of reaction conditions. In order to explore this issue in more detail, the release of HC isomers from a Zn/Cu HDSs via exchange with chloride anions has been investigated. The overall reactivity was assessed at 40° C. After exposure to 1 M aqueous solutions of NaCl, the extent of reaction at equilibrium was assessed by characterizing the amount of HC isomer released into solution. Under these conditions, 90% of o-HC was found to be released via the exchange reaction while only 40% was released from m-HC and 22% from p-HC. The anion with the largest dipole moment, p-HC, exhibited the lowest extent of reaction while the anion with the lowest dipole moment, o-HC was released nearly to completion under the same conditions.

The kinetics of these reactions can be evaluated via model-based or model-free approaches, as described in detail in recent work from our laboratory [12]. The Avarami-Erofe'ev model is commonly applied to these types of reactions. At 40° C, the rate constants obtained via Avarami-Erofe'ev analysis were 7± 2 xlO"3 s"1 for m-HC, 1.0 ± 0.2 xlO"3 s'1 forp-HC, and 0.29 ± 0.05 xlO"3 s"1 for o-HC. The rate of release follows the following order. m-HCn >p-HCn > o-HCn. The kinetics therefore do not follow the same trend as the equilibrium extent of reaction.

A key factor in the trends in anion release kinetics appears to be the orientation and extent of hydrogen bonding of the anions in between the HDS layers. Figure 2 shows ATR-FTIR analysis of the hydroxy cinnamate HDS structures, along with the same data for the Zn/Cu acetate precursor that was used in the synthesis of the hydroxy cinnamate HDSs.

Figure 1 : FTIR spectra for Zn/Cu hydroxy cinnamate structures.

The observed d-spacings for the hydroxy cinnamate HDSs are 16.95,20.72, and 21.03 Â for the ortho, meta, and para isomers, respectively. The hydroxy cinnamate chain lengths can be

408

estimated as approximately 8.6 Â for the ortho and meta isomers and 9.1 Â for the para structure. Taking into account the thickness of the metal hydroxide layers, this suggests a slightly tilted bilayer orientation for the meta and para structures. The smaller d-spacing observed with ortho hydroxy cinnamate is consistent with a more tilted orientation within the interlayer gallery.

The FTIR data provide evidence for differences in the extent of hydrogen bonding via inspection of the hydroxyl stretching region (~ 3000-3500 cm"1). The broader peaks observed in the ortho case are consistent with more extensive hydrogen bonding in the case of the ortho isomer. Combining the XRD and FTIR data, the interlayer orientation of the meta isomer is proposed to tilt the anion toward the metal hydroxide layer in order to enable hydrogen bonding between the hydroxyl group and the layer. This, in turn, is the cause of the smaller-than-expected d-spacing when the ortho -containing HDS is compared to the meto-containing structure. The observed rate constants are inversely proportional to the extent of hydrogen bonding. The ortho isomer exhibits the slowest release. The para also exhibits some hydrogen bonding which we assign to tail-to-tail interaction between hydroxyl groups on two hydroxy cinnamate molecules.

When designing layered metal hydroxide species for controlled release applications, one therefore must consider at least three factors. The metal ion composition of the metal hydroxide layer plays a role in determining the relative ease of releasing the interlayer anion via a simple exchange reaction. Further work is necessary to determine the factors involved in the metal ion effect on reactivity. In the case of the hydroxy cinnamate model system, the anion dipole moment is inversely correlated with the extent of reaction (i.e. the thermodynamics) while hydrogen-bonding interaction of the anions with the metal hydroxide layer or with other interlayer anions slows the rate of release.

General Reactivity Trends: Nitrate/Acetate Exchange Reactions

Another potential application for layered metal hydroxides is use for removal of aqueous contaminants, either by adsorption or by anion exchange. A targeted application in our laboratory is the removal of aqueous contaminates, such as chlorinated acetates. As a model system, we are exploring the uptake of acetates by nitrate-containing layered metal hydroxides, as well as the reverse reaction, release of acetate by nitrate exchange. An overview of the general reactivity trends for Zn-based layered metal hydroxides is provided in Table II for reactions performed at 40° C. In the case of release of nitrate by exchange with acetate, the copper-containing HDS (ZnCuN) does not react, while the other three materials do exchange. Copper hydroxyl nitrate also does not exchange with acetate. Preliminary analysis of FTIR data suggests that in the case of the copper-containing materials, the nitrate is bound to the metal hydroxide layer since the nitrate bands are consistent with C2V symmetry that is expected for a grafted structure. The other nitrates listed in Table II exhibit FTIR spectra consistent with unbound (D3h symmetry) nitrate. Therefore the ease of nitrate release is optimized for unbound structures. Further work is in progress to characterize other hydrogen bonding effects and the potential role of interlayer water.

409

Table II: General Reactivity Trends in Acetate/Nitrate Exchanges

Nitrate precursor

ZnHN Zn/CuHN Zn/NiHN Zn/CoHN

Exchange with acetate?

yes no yes yes

Acetate Precursor

ZnHAc Zn/CuHAc Zn/NiHAc Zn/CoHAc

Exchange with nitrate?

yes yes no no

The reverse reaction, release of acetate via exchange with nitrate, occurs for Zn hydroxy acetate (ZnHAc) and the Zn/Cu HDS (Zn/CuAc). The forward and the reverse reactions are observed only in the case of the Zn metal hydroxides. Figure 3 shows the results of monitoring nitrate concentations in solution for die uptake (i.e. the reaction ZnHAc + nitrate) and the release (ZnHN + acetate) reactions.

Figure 3. Solution phase nitrate concentrations (both runs performed under identical conditions).

410

Conclusions

The nature of the anion interlayer interactions (bonding and hydrogen bonding) with the metal hydroxide layer or with other interlayer anions is an important factor in determining exchange reactivity. Applications that involve either incorporation of a new anion into the interlayer domain or controlled release of the anion require consideration of nature of these interactions. The metal ion composition also influences reactivity; further work is necessary to understand the controlling factors.

Acknowledgments

This work was supported by the National Science Foundation (CHE-0809751 ).

References

1. V. Rives, (ed.) Layered Double Hydroxides: Present and Future (New York, NY: Nova, 2001).

2. W.T. Reichle, "Catalytic reactions by thermally activated, synthetic, anionic clay minerals," J. Catal., 94 (1985), 547-557.

3. C. Kappenstein, et al., "Copper-zinc oxide catalysts. Part IV. Thermal treatment in air, argon and hydrogen and XRD study of new bimetallic precursors-direct formation of alloys, " ThermochimicaActa, 279 (1996), 65-76.

4. J. Cemak, et al., "Copper-zinc oxide catalyst. Part II. Preparation, IR characterization and thermal properties of novel bimetallic precursors," ThermochimicaActa, 276 (1996), 209-220.

5. Y. Guo, R. Weiss, M. Epple, "A Straightforward Route to Copper/Zinc Oxide Nanocomposites: The Controlled Thermolysis of Zn[Cu(CN)3]," Eur. J. Inorg. Chem., 2005, 3072-9.

6. E.L. Crepaldi, P.C. Pavan, J.B. Valim, "Anion exchange in layered double hydroxides by surfactant salt formation," J. Mater. Chem., 10(2000), 1337-1343.

7. C.S. Bruschini, M.J. Hudson, M. J., "Anion exchange in copper hydroxy double salts," Special Publication - Royal Society of Chemistry (Progress in Ion Exchange), 196 (1997), 403-411.

8. M. Meyn, K. Beneke, G. Lagaly, "Anion-exchange reactions of hydroxy double salts," Inorg. Chem., 32 (1993), 1209-1215.

9. (a) M. Bellotto, et al., "A Reexaminataion of Hydrotalcite Crystal Chemistry," J. Phys. Chem. 100 (1996), 8527-34. (b) B. Mavis, M. Akinc, "Cyanate Intercalation in Nickel Hydroxide," Chem. Mater., 18 (2006), 5317-25.

10. (a) J. Choy, et al., "Structural phase transformation of layered hydroxyl double salts, Nil-xZn2x(OH)2(CH3COO)2xnH20, depending on hydration degree," Bull. Korean Chem.

411

Soc, 18 (1997), 450-3. (b) H. Morioka, et al., "Effects of Zinc on the New Preparation Method ofHydroxy Double Salts," Inorg. Chem., 38 (1999), 4211-16.

11. (a) W. Fujita, K. Awaga, "Reversible structural transformation and drastic magnetic change in a copper hydroxides intercalation compound," J. Am. Chem. Soc, 119 (1997), 4563-64. (b) V. Läget, et al., "Hybrid organic-inorganic layered compounds prepared by anion exchange reaction: correlation between structure and magnetic properties," J. Mater. Chem., 9 (1999), 169-174. (c) M. Taibi, et al., "Layered nickel hydroxide salts: synthesis, characterization and magnetic behavior in relation to basal spacing," J. Mater. Chem., 12 (2002), 3238-44.

12. S. Majoni, J.M. Hossenlopp, "Anion Exchange Kinetics of Nanodimensional Layered Metal Hydroxides: Useoflsoconversional Analysis,"./. Phys. Chem. A, 114 (2010) 12858-12869.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ELECTROCHEMICAL SYNTHESIS OF LAYER DOUBLE HYDROXIDES, ITS CHARACTERIZATION, AND PERFORMANCE

STUDY FOR REMOVAL OF NITRATE AND ARSENIC

Md. Iftikher Haider1, Jewel Andrew Gomes1, Kevin Urbanczyk2, David L. Cocke3, Hylton McWhinney4, George Irwin5, Paul Bernazzani6

'Dan F. Smith Department of Chemical Engineering, Lamar University, Beaumont, TX 77710 department of Geology, Sur Ross State University, Alpine, TX 79832

3Gill Chair of Chemical Engineering, Lamar University, Beaumont, TX 77710 4Department of Chemistry, Prairie View A&M University, Prairie View, TX 77446

'Department of Physics, Lamar University, Beaumont, TX 77710 department of Chemistry and Biochemistry, Lamar University, Beaumont, TX 77710

Keywords: Electrocoagulation, green rust, iron, aluminum, magnesium, sulfate

Abstract

Green rust (GR) is an important intermediate Layered Double Hydroxide (LDH), which exists in oxidative transformation of iron(II) phase. Anionic and cationic species can replace their corresponding parts in the layered structure, and different species can also be accumulated in between the layers. GR can play an important role in wastewater remediation, such as reduction and removal of toxic compounds. An afford was carried out to synthesize LDH electrochemically, and to characterize it by Mössbauer, x-ray diffraction, scanning electron microscopy/energy dispersive x-ray analysis, x-ray photoelectron spectroscopy, and Fourier transform infrared spectroscopic techniques. GR was used to remove nitrate and arsenic from wastewater. The optimized conditions for removal of these water pollutants were also investigated. The resistance against oxidation of GR was increased using dopents, such as magnesium and aluminum, so that it can be used for water treatment in remote areas.

Introduction

Green rust (GR) is a unique class of layered double-hydroxides (LDH's) that is represented by the general formula: [Mlf

(i_x) M111

x(OH)2]x+ [(x/n) An\ mH20]x" where both M11 and M111 are

metallic cations , x= Feln/Fet01al, An" denotes interlayer anions whose charge is n, and m is moles of water molecules. Green rusts (GRs) had been studied for two major purposes. First, they are intermediate products during the electrocoagulation process and consequently play a role in the aqueous corrosion of iron based materials. Second, they have been identified as a mineral of bluish-green color in anaerobic conditions that turns ochre once exposed to the air [1]. The crystal structure of GR consists of positively charged brucite-like layers of edge-sharing octahedrally coordinated Fe" and Fe1" hydroxide units intercalated with anionic species and water molecules [2], GR are interesting compounds due to the presence of the reactive Fe" cations that can easily reduce anionic pollutants such as nitrate, selenate or chromate. The overall spectra of LDHs have

413

received increasing interest in recent years owing to their applications as catalysts, adsorbents and ion exchangers [3]. Formation of LDHs is governed by many parameters such as the nature of the cations, the pH of precipitation, temperature and the precipitation method. Physical parameters that can be controlled are the crystal size, morphology and specific surface area (SSA) for optimizing the LDH reactivity [4]. Electrocoagulation produces green rust as intermediates, and it is responsible for removing water contaminants. Therefore, the target of our study was to treat wastewater without in-situ electrochemical experiment. To meet our goals, we obtained optimized conditions for electrochemically generation of green rust (GR), and observed stability in aerobic condition by inclusion of aluminum and magnesium in the green rust structure. As part of characterization, samples were analyzed by Moessbauer, XRD, SEM/EDS, XPS, TG/DSC and FTIR. At the final stage performance study was performed for both original and modified green rust for arsenic and nitrate removal in wastewater.

Experimental

Batch electrocoagulation experiments were carried out in a beaker with 400 mL solution using vertically positioned Fe-Fe electrodes dipped in the solution. The anode and cathode were with dimensions of 3.2 cm * 6.2 cm, both electrodes having a submerged surface area of 14.4 cm2. The electrode spacing was 2.5 cm. Electrodes were connected to a DC power supply (Kaselco). Electrocoagulation experiments were performed at a constant current of 1 A while potential varies from 8-50 V depending on the varied conductivity. Prior to each experiment, the electrodes were abraded with sand-paper to remove scale, then treated with a solution of acetone in order to reject any effect due to the different prehistory of the electrodes, washed with distilled water, dried and then, the anode was weighed. An appropriate amount of HC1 (37%) or NaOH (1.125 M) added at the beginning of the process and during generation of GR formation. Moreover, pH was measured continuously using a Denver Instrument UB-10 pH meter. Electrocoagulation experiments were run for 10 minute. When GR produced, the beaker containing GR was sealed immediately and ultra pure nitrogen Gas (purity 99.9999 %) was flowing continuously for 15 min. to remove the dissolve oxygen to ensure that no further oxidation occurs. After that the solution was quickly transferred in a glove bag and the bag was purged with nitrogen gas for several times in order to avoid the oxidation of ferrous ions (Fe(ll)). On that condition, the filtration was performed. Floes were separated by a filter paper and kept in glove bag for 12-15 hours to dry completely under regular air purging from the glove bag. Then the sample was collected in a sealed box and paste with tape or para-film and referred for further characterization.

Results and discussion

Optimized pH for GR Formation

The experiment was done in the pH range 2-10. From the results it can be concluded that if the pH is too low (less than 3.98), Fe(II) would be oxidized directly into Fe (III) precipitates. Besides, if the pH is basic in excess (greater than 10.2), magnetite, Fe3Û4, would be formed.

414

From this observation, it can be concluded that the nucleation of GR start at slightly alkaline condition (greater than 7.2) and the formation continue until 10.2. Besides, when GR formation start, if acid is added, GR turned into yellowish product. That indicates that GR is stable only in slightly alkaline condition. Therefore, there was generation of ferric species (oxidized product of GR) at lower pH, and magnetite at higher pH.

Impregnation of GR with Mg" and / or Al1"

The objective of this work was to investigate the influence of the extent of M(lll) substitution in LDHs. Three series Fe-Al, Mg-Al-Fe and Mg-Fe wer e synthesized with different Mg:Al:Fe ratio and x in the range 0.15 - 0.7. Using Faraday's second law of electrolysis, the amount of iron dissolved was calculated. Based on this amount, definite amount of magnesium and aluminum salts were added so that the mole ratio among these three metal ions become the number as desired. The influences of the nature and the extent of M(III) ion substitution on the structure, morphology and surface properties of layered double hydroxides, and derived mixed oxides were investigated.

Material Characterization: FTIR Two samples (Bl and B2) were prepared by (15:83) and (83:25) molar ratio of Al and Mg, respectively. The FTIR spectrum shown in Table I clearly indicates that the band frequencies of GR-SO4 are almost similar to the value that were found in literature for GR prepared by co-precipitate method [2], Sample Bl showed more pure GR characteristic, whereas B2 showed some oxidized characteristics. Therefore, it can be concluded that stability of GR-SO4 had enhanced through increased doping of magnesium.

Table I. Comparative spectrum data for sample B1 and B2

Bl

760

1,098

1,630

3,430

B2

-1075

1622

-

Pure GR [2]

773 B1

1,102 Bl

1,657 B'

3,469B1

Oxidized GR [2]

-1060 B2

1644 B2'

3,070

Tentative assignment

Fe-OH Modes

Spitted V3.SO4

8-H20

v-OH

Material Characterization: XRD The sample was prepared by impregnating (28 mole %) Al(III) and (43 mole %) Mg(II). The remaining percentage amount was iron. X-ray diffraction was performed using D4 Endeavor Bruker instrument (Figure 2, black pattern). It was observed that the sample contains mostly iron hydroxide sulphate green rust, hematite and maghemite. The sample was exposed in air to have oxidation for 20 min., and XRD was run again on it (Figure 2, red pattern). It was found that the intensity of the diffraction peaks green rust was diminished or significantly reduce due to oxidation. On the other hand, no change in the intensity was found for the diffraction peaks for solely hematite and maghemite. The result indicates that the addition of aluminum and magnesium increase the stability of green rust. In addition, inclusion of Al and Mg produces an increase of amorphousity or nano-crystallinity, and thus possibly an increase in surface area

415

Material Characterization: Thermal analysis Four samples were prepared with the following composition: Jl (Al:Mg:Fe = 0:59:41), J2 (Al:Mg:Fe = 49:0:51), J3 (Al:Mg:Fe = 28:42:30), and H4 (Al:Mg:Fe = 51:29:20). Thermal analyses were made in helium atmosphere at the temperature ramp of 10 °C/min up to 800 °C. TG curve (percent wt loss vs. temperature) and DSC curve (heat flow variation with temperature) were shown in Figure 3 for Jl sample.

Figure 1. XRD patterns of green rust (black) and oxidized green rust (red).

Figure 2. % Wt Loss and Heat Flow (mW) variation vs Temp (CC) for sample Jl.

For every sample, two endothermic and one exothermic region were found. The temperature ranges of the endothermic peaks and the corresponding weight loss were tabulated in Table II. The first endothermic region was observed starting from 73, 85 107, and 107 °C for the sample Jl, J3, J2, and H4, respectively. This region is dedicated to two differently bonded water

416

molecules release during the heating process. First, loosely bonded water present in the sample and second one is for the crystal bonded water. The region was continued by the dehydroxylation of the hydroxyl groups bonded with different layer cations. In particular, as bond strength of Mg-O (375.8 kJ/mol) is lower than Fe-O(407.2 kj/mol), the deoxylyation of Mg should start before Fell/FelH oxide [5].The dehydroxylation of the hydroxyl groups bonded with of Mg and Fe was taken place at the temp of 126 and 151 °C, respectively.

Table II. Temparature ranges of the endothermic peaks found for the modified GR samples.

Sample ID

Jl J3 J2 H4

Endo-1 Peak Range in °C

Start

73 85 107 107

Stop

208 218 274 240

wt loss, %

8 7 9 10

Endo-2 Peak Range in °C

Start

218 242 249 239

Stop

258 291 297 272

wt loss,%

2 2 2 2

In case of sample Jl (0% Al), there is 8% wt. loss in this region. But with Al doping, wt loss had decreased by 1% for sample J3 (28% Al). But for sample J2 (49% Al) the wt loss was increased to 9%, and for the sample H4 (51% Al), it is 10%. In addition, the starting temperature of this peak was increased from 73 °C to 107 °C with the increase in addition of aluminum. The result indicated that bonding strength of water in GR had been increased significantly with certain amount of AI doping. For sample J2, the range between the start and end temperature was maximum, i.e., 167 °C. But when impregnation of magnesium was involved, there was no change in the difference of temperature. It means that the thermal decomposition in this region was retarded or reacted less sensibly by the inclusion of magnesium. For the second endothermic region, when there is no aluminum (sample Jl), the wt loss before second endothermic region is 9.5%. By doping of aluminum, wt loss was decreased to 8.8% for J3, and 9.0% for J2. But after the certain amount of impregnation of aluminum, the amount was increased from 9.51% to 10.85%. Therefore, it can be concluded that impregnation of aluminum up to certain number of moles might contribute to the thermal stability of the green rust.

Bonding strength of Al-O (514 kJ/mol) is higher than Mgll and Fell/Ill. The second endothermic peak is attributed to the dehydroxylation of hydroxyl groups bonded with Fe III followed by Al III.

Zhang et al. [6] stated that there was removal of interlayer SO42" anions around 338 °C. So the exothermic peak of our experiment (peak maxima are in the range of 343 to 386 °C for different samples) was attributed to the slow mass loss event by removal of small amount of interlayer SO4 anions, followed by decomposition SO42' present in the GR structure.

If we consider the weight loss up to 793 °C, we found that the maximum wt loss took place for sample Jl where there was no aluminum. When aluminum was doped in the GR structure, the weight loss is reduced. Such as, for sample J2 and J3, it is about 5%, and for H4, it is 8%.

417

Therefore, it can be concluded that aluminum inclusion retards the thermal decomposition of GR after removal of interlayer SO42" anions.

SEM/EDAX SEM images indicate that GR-SO4 layer is very compact and exhibits thick platy particles. In addition, it has a different morphology displaying a honeycomb structure, typical of some double layered-hydroxides compounds, with fine and intermingled particles. EDAX spectra confirm the presence of Al/Mg, Fe, O, and S in the GR-SO4 samples.

Moessbauer From Moessbauer spectra of GR showed quadrupole doublets corresponding to the Fell and Felll in the material. However, only a narrow quadrupole doublet corresponding to Fe III in the decomposition products were observed, characteristic of fine particles of iron oxyhydroxides: goethite or lepidocrocite.

Performance studies of green rust

Performance Study for Arsenic Removal. The prepared GR including the modified ones were used to treat arsenic containing synthetic waters. To find out the arsenic removal efficiency of green rust, arsenic solution of 60 ppm was mixed with GRs and kept overnight. It has found that the amount of green rust prepared by 15 min EC run provides 96.6% removal efficiency of arsenic, whereas for the 10 min EC run, the arsenic removal efficiency was found as 96.4% that is more economical considering the power consumption.

Arsenic Removal by Mg 2+ Impregnated GR. In the next step, Mg2+ was impregnated in GR in different proportion and was used to treat arsenic. It has found that 54.3 molar % inclusion of magnesium gives 97.75% removal efficiency.

Arsenic removal by Al3+ impregnated GR. In the next step, Al3+ was impregnated in GR in different proportion and was used to treat arsenic contaminated water. It had found that simulation-wise it is possible to achieve removal efficiency up to 100% with the aluminum doped GR (Figure 4).

Figure 3. Arsenic removal with Al + containing green rust. The black colored curve indicates the theoretical trends of arsenic removal efficiency prepared by MATLAB.

418

Arsenic removal by both Mg2+ and Al3+ impregnated GR. The maximum removal efficiency 98.6 % was found by the GR doped with both aluminum and magnesium.

Figure 4. % Arsenic removal with both Al3+ and Mg2+ combined green rust.

Performance study by nitrate removal. The performance of the prepared GR was tested by applying it in a aqueous solution containing potassium nitrate of 37.8 mg/L. The data was continuously measured by an ion selective electrode (Vernier) for first four hours (Figure 6). Then it was left for overnight removal. After overnight application of the GR, the solution was filtered and the nitrate concentration was measured again by ISE. The concentration after treatment was found 13.5 mg/L.

Figure 5. Measured data of nitrate concentration at first four hours application time of

green rust.

From the result it was found that 64.3% nitrate can be removed by pure green rust ex situ after overnight application. According to Christian et al. [7] and Hansen et al. [8] who prepared GR chemically, if nitrate can exchange with the interlayer anion (A"~), reaction can take place both at outer and inner surfaces of the GR particles. The rate of reaction will depend on the particular anionic form of GR, specifically the ease with which nitrate can exchange with A"" in the GR interlayer, and the crystallite size. It also depends on the layer charge and the relative content of Fe(Il) in the hydroxide layers. Further investigation is required to illustrate the nitrate reduction

419

mechanism with the electrochemically prepared GR.

Conclusion

Electrocoagulation technique can be efficiently used to prepare pure and modified green rust electrochemically. In the synthesis process, nucleation starts at slightly alkaline condition (pH greater than 7.2), and the formation continues until 10.2. Other pH conditions lead to different oxidative products. Inclusion of aluminum and magnesium cations in the GR structure increases its stability. XRD analysis confirms the formation of green rust and some other oxidized products, such as iron oxides and oxyhydroxides in the samples. It also indicates that Al and/or Mg doping increases amorphousity of green rust. FT1R analysis confirms the presence of Fe-OH, SO4, and H20 in the prepared GR. SEM analysis shows some shapeless amorphous morphology of green rust.TG-DSC results refer to the endothermic processes of dehydration (loose and crystalline H2O), dehydroxylation (cation bonded OH"), and the exothermic process of SO4 removal. Modified green rust doped with aluminum and magnesium provides the best arsenic removal efficiency (98.6%). Pure GR provides 64.3% of nitrate removal efficiency.

Acknowledgements

We greatly acknowledge the financial support from USDA (2009-38899-20017) and Lamar Research Enhancement Grant (2009), and the instrumental support from Lamar Materials Instrumentation Center.

References

1. L, Legrand., M. Abdelmoula, A. Genin, and A. Chausse. "Electrochemical formation of a new Fe(II) /Fe (III) hydroxy- carbonate green rust: characterization and morphology." Electrochimica Ada 46 (2001), 1815-1822. 2.1.A,M. Ahmed et al, "Formation of Green Rust Sulfate: A Combined in Situ Time-Resolved X-ray Scattering and Electrochemical Study," American Chemical Society, 26 (9) (2010), 6593-6603. 3. C. Ruby et al., "Synthesis and transformation of iron-based layered double hydroxides," Applied Clay Science, 48 (1-2) (2010), 195-202. 4. C. Ruby et al., "Oxidation modes and thermodynamics of Fell—III oxyhydroxycarbonate green rust: Dissolution-precipitation versus in situ deprotonation," Geochimica et Cosmochimica Ada 74 (2010), 953-966. 5. D.R. Lide, edi., CRC Handbook of Physics and Chemistry, 89th Edition (New York: CRC Press, 2008). 6. H. Zhang et al., "Syntehsis and characterization of sulphate and dodecylbenjenesulfonate intercalated zinc-iron layered double hydroxide by one step co-precipitation route," Solid state chem 180 (2007), 1636-1647. 7. H. Christian et al., "Kinetics of nitrate reduction by green rusts—effects of interlayer anion and Fe(II) / :Fe(III)/ ratio," Applied Clay Science 18 (2001), 81-91. 8. B., Hansen et al., "Reduction of nitrate to ammonium by sulphate green rust: activation energy and reaction mechanism," Clay Minerals, 33 (1) (1998), 87-101.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

REMOVAL OF DIRECT RED AND ORANGE II AZO DYE FROM SYNTHETIC TEXTILE WATER USING ELECTROCHEMICALLY

PRODUCED FE-LDH

Sadia Jame1, Jewel Andrew Gomes2, David L. Cocke3

'Department of Civil Engineering, Lamar University, Beaumont, TX 77710 2Dan F. Smith Department of Chemical Engineering, Lamar University, Beaumont, TX 77710

3Gill Chair of Chemical Engineering, Lamar University, Beaumont, TX 77710

Keywords: Direct Red, Orange II, Electrocoagulaton, Iron

Abstract

Layered double hydroxides are produced as an intermediate phase during in-situ electrochemical coagulation technique using iron sacrificial electrodes. Its structure consists of cationic layers, Fe(OH)2n+, alternating with anionic layers stacked in a six-layer repeating sequence, giving rise to a hexagonal or pseudohexagonal symmetry. This Fe-LDH has the ability to trap organic impurities in the interlayers when it is mixed with wastewater. This hypothesis was verified using dye molecules, such as Direct Red and Orange II. Generally, textile industries use various kinds of chemicals including dyestuffs and pigments during dyeing and finishing processes. Their effluent is extremely toxic and environmentally hazardous. Fe-LDH was found to remove dye molecules with satisfactory efficiency. Optimum conditions were explored by changing current density, dye concentration, conductivity and pH during the LDH generation process. The floe was characterized using x-ray diffraction.

Introduction

Layered double hydroxides (LDHs), also known as hydrotalcite-like compounds or anionic clays, have received much attention in the past decades due to their vast applicability. An LDH is created when a fraction of the divalent cations in a brucite-like lattice are isomorphously replaced by trivalent cations, introducing a positive charge in the layers. This charge is electrically balanced by anions located in the interlayer region, along with hydration water molecules. Given the wide range of compounds with LDH structure that may be prepared, they are represented by the general formula: [M"(|. X(M"x (OHM A""^ -mHiO, where M1 includes: Mg2+, Co2+, Cu2+, Ni2+, Zn2+, etc.; M1" may be Al3+, Cf+, Ga3+, Fe3+; and A"" might be any organic and/or inorganic anion. Many ternary LDHs involving mixtures of different M11 and/or M ' may also be prepared [1, 2, 3], LDHs are also known for their ability to adsorb negatively charged species and/or pollutants, which is mainly due to their large surface area, high anion-exchange capacities and flexible interlayer space. It has been successfully tested in the removal of pollutants, such as the herbicide 2,4-dichlorophenoxiacetic acid (2,4-D), from aqueous solutions [4]. Additionally, this compound was used as an adsorbent for removal of color from synthetic textile wastewater (STW) and textile wastewater (TWW) [5]. Wastewaters generated by textile industries contain large amounts of toxic aromatic compounds, especially azo dyes. Azo dyes and their degradation products, such as aromatic amines, are highly carcinogenic. These effluents are also responsible for causing high levels of color. The colored wastewater

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released into the environment is also a source for disruption of aquatic life and aesthetic pollution as well [6]. Adsorption and precipitation processes are very time-consuming and as well as costly with low efficiency. Chemical degradation is one of the most important and effective methods, but it produces some very toxic products such as organo-chlorine compounds [7], In the recent years, investigations have been focused on the treatment of wastewaters using electrocoagulation (EC) because of the increase in environmental restrictions on effluent wastewater. Indeed, EC is a simple and efficient method for the treatment of most drinking waters and wastewaters as well. EC as an electrochemical method was developed to overcome the drawbacks of conventional water and wastewater treatment technologies. It provides a simple, trustworthy and cost-effective method for the treatment of wastewater without any need for additional chemicals. EC also reduces the amount of sludge, which needs to be disposed [8]. Iron and aluminum are the most widely used electrode materials in EC process. It has been tested successfully in the separation of pollutants from restaurant wastewater, treatment of urban wastewater [9], degradation and decolorization of dye solution [10-11], defluoridation of water [12-13], separation of aqueous suspensions of ultrafine particles [14], and removal of nitrate from water [15]. Therefore, it is expected that EC would be an ideal choice for decolorization of dye solutions [16]. However, many of the commercially available azodyes show mutagenic activity in vitro due to impurities, e.g. contamination with aromatic amines. The objective of the present study is to investigate the decolorization of mixed dye of Direct Red 23 and Orange II in aqueous solution using iron electrodes along with floe characterization.

Experimental

Chemicals Azo dyes: Direct Red 23 and Orange II dye were used in this work. Mass ratio of two dyes was held constant at 1:1. Dyes were used for preparing synthetic waste water by dissolving these in distilled water. In addition, sodium chloride (NaCl) was used to increase the conductivity of the solution, and granular sodium hydroxide and/or sulfuric acid were used to maintain the pH. The structures of two dyes are shown in Table I.

Apparatus The beaker size EC was carried out in a 250 mL beaker with magnetic stirrer, using vertically positioned electrodes dipped in the wastewater. Two iron electrodes were used with the effective dimension of 3.8 cm x 3.0 cm x 0.1 cm. The current and voltage during the EC process were checked using a Cen-Tech multimeter. The pH of the solutions before and after EC was measured by an Oakton pH meter. EC was run for a certain period of time according to the experiment. The experiments were performed at room temperature of 25 °C. After EC-run, the EC-mixture was filtered and the precipitate was dried.

Analytical method A UV-spectrophotometer (Shimadzu, Model UV-160) was employed to measure the absorbance of the dye solutions at maximum wavelength (AmiK = 493 nm) that was used to measure the concentration of the dye. This absorbance band was found to be common for both the dyes. The decolorization efficiency, E, is calculated as

422

£=(Ci-Cf)/CiXlOO (1)

Where, Ci is initial dye concentration (mg/L) and Cf is final dye concentration (mg/L). The XRD analysis of the Electrocoagulation by-products were carried out with a Bruker D8 GADDS Discover diffractometer operating with Cu K„ radiation source filtered with a graphitic monochromator.

Table I. Structures of the dyes: Direct Red and Orange II.

Results and discussion

Decolonzation efficiency was investigated at different conductivity, pH and current density and dye concentrations. Both the dyes have significant absorbance peaks at 493 nm, and we measured the concentration change of the dyes by measuring the absorbance at this wave-length with EC run for 16 minutes.

Effect of Electrolyte Concentration The amount of sodium chloride (NaCl) was varied from 2 g/L to 5 g/L liter in the reactant solutions. EC was run for each of these solutions for 16 minutes at current density of 248 A/m2. The results are shown in Figure 1. The highest removal efficiency (RE %) of 98% was found at 3 g/L of electrolyte concentration. A similar kind of effect of increase in removal efficiency with increase conductivity was also reported by Kashefialasl and others [17].

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Figure 1. Effect of electrolyte concentrations.

Effect of Initial pH pH is an important operation factor influencing electrocoagulation process. The effect of initial pH with Fe sacrificial electrodes is shown in Figure 2.

Figure 2.Effect of Initial pH.

The %RE versus pH clearly indicates that the removal efficiency increases slowly with pH and it remains unchanged between pH 6.0-8.0. Maximum (94%) and minimum (74%) removal efficiencies occurred at 7 and 10 [18].

Effect of Current Density Figure 3 shows the removal efficiency as a function of current density. As the applied current density was increased from 83 to 414 A/m2, 248 A/m2 can be considered realistic optimum current density with good removal efficiency (96%) considering the cost for electric power.

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Figure 3. Effect of Current Density

Effect of Dve Concentration Five experiments with different concentrations of 9 ,11,13 and 15 ppm of mixed dye solution with current density 155 A/m2, electrolyte concentration 3 g/L were carried out. The mass ratio of two dyes was held constant at 1:1. The maximum removal efficiency with optimum concentrations was determined. As shown in Figure 4, the removal efficiency decreases linearly from 88% to 71% as the concentration increases from 10 ppm to 15 ppm. This might happen due to insufficient generation of floe during 10 minutes of EC run for each case. From our results, the optimum concentration of the mixed dye solution was found to be 10 ppm.

Figure 4. Effect of Dye Concentration.

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Optimum Conditions Among the ranges of variables we performed the EC experiments, the optimum conditions were found as follows: current density 248 A/m2; initial pH 7.0; dye concentration 10 ppm; and electrolyte concentration 3 g/L.

Material Analysis X-Rav Diffraction. Figure 5 shows the XRD results for the characterization of EC floes. The EC floe shows the presence of magnetite (cubic), iron oxide (tetragonal) and goethite (orthorhombic) produced during EC reactions. The absence of mixed dye peaks in the EC-floc possibly indicate the degradation of DR 23 molecules and Orange II during EC.

Proposed Mechanism It has been observed that during Electrocoagulation (EC) process, a green color floe was produced by dissolving metal ions that in turn removes the pollutants from water. According to Tamura [19] green rust (GR) may produce as follows:

2Fe'"(OH)2+ + 4Fe" OH+ + 2CF + 40H" + 3H20 «- [F^Fe"1

2(OH) i2] [Cl2" *H20] (2)

Both Direct Red 23 (DR) and Orange II dye are ionic in water. They may be dissolved as follows:

Dye-NaS03 Na++ [Dye-S03]" (3)

Anionic part of dissolved dye molecule may electrostatically adsorb on the GR surface by the following manner:

[Fe'^Fe1112(0H),2][C12" -H20]+ 2[Dye-S03]" ->■ [Fe'^Fe111

2(0H),2] 2[Dye-S03] +C12+ 2e + 3H20 (4)

During the chemical transformation of GR to iron oxides, the anionic dye part is probably degraded to other species. As mentioned by Venarkova et al. [20], aromatic degradation products can be obtained due to the cleavage of azo bonds and chlorination of the products are also possible. XRD analysis of the EC floes indicated presence of two compounds namely, iron oxide, goethite (FeOOH), and magnetite (Fe30<t). The oxidation of the Fe2+ ions in GR results in Fe203-Fe304 in oxygen- depleted systems through FeOOH. The formation of magnetite in the by-products was discussed by Satapanajaru et al. [20].

Conclusion

Electrocoagulaton with iron sacrificial electrodes is very efficient to remove azo dyes, such as Direct Red 23 and Orange II. The optimum conditions for maximum removal of dyes (96%) found in this work are as follows: current density 248 A/m2 ; initial pH 7.0; dye concentration 10 ppm; and electrolyte concentration 3 g/L. Tetragonal iron oxide, orthorhombic geothitte, and cubic magnetite were found in the EC-floc confirming the formation of green rust as intermediates. The mechanism of removal of azo dyes in EC were also illustrated.

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Figure 5. XRD of Direct Red 23 and Orange II mixed dye (black), and EC-floc (sky-blue).

Acknowledgements

We greatly acknowledge the instrumental support from Lamar Materials Instrumentation Center.

References 1. Y. Song et al., "Self-assembled hierarchical porous layered double hydroxides by solvothermal method and their application for capacitors," Microporons and Mesoporous Material, 148 (1) (2012),159-165.

2. S.Morandi et al., "Supported Ni catalysts prepared by intercalation of Layered Double Hydroxides: Investigation of acid-base properties and nature of Ni phases," Microporous and Mesoporous Materials, 147 (1) (2012), 178-187.

3. D.E. Giles et al., "Iron and aluminium based adsorption strategies for removing arsenic from water," Journal of Environmental Management, 92 (12) (2011), 3011-3022.

4. N. Iyi, H. Yamada, and T. Sasaki, "Deintercalation of carbonate ions from carbonate-type layered double hydroxides (LDHs) using acid-alcohol mixed solutions," Applied Clay Science, 54 (2) (2011), 132-137.

5. M.Z. Hussein et al., "The use of Mg/Al layered double hydroxide for color removal of textile wastewater," Journal of Environmental Science Health A Toxic/ Hazard Substance Environmental Engineering, 36(4) (2001), 565-73.

6. M.N. Pons et al., "Electro-coagulation of reactive textile dyes and textile wastewater,' Chemical Engineering Process, 44 (2005), 461^170.

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7. T. Kim et al., "Decolorization of disperse and reactive dyes by continuous electrocoagulation process," Journal of Desalination, 150(2002), 165-175.

8. M. Y.A. Mollah et al., "Fundamentals, Present and Future Perspectives of Electrocoagulation," Journal of Hazardous Materials, Bl 14 (2004), 199-210.

9. G. Chen, "Electrochemical technologies in wastewater treatment," Separation and Purification Technology, 38 (1) (2004), 11- 41.

10. Z. Shen, et al., "Degradation of dye solution by an activated carbon fiber electrode electrolysis system," Journal of Hazardous Material,84(B) (2001), 107- 116.

11. J.S. Do, and M.L. Chen, "In situ oxidative degradation of formaldehyde with hydrogen peroxide electrogenerated on the modified graphites," Journal of Applied Electrochemistry, 24 (1994), 936-942.

12. N. Mameri et al., "Deflouridation of Sahara water by small plant electrocoagulation using bypolar aluminum electrode," Separation and Purification Technology, 24 (2001), 113-119.

13.N. Mameri et al., "Defluoridation of septentrional Sahara water of North Africa by electrocoagulation process using bipolar aluminum electrodes," Water Research, 32 (50) (1998), 1604-1612.

14. M.J. Matteson, et al., "Electrocoagulation and separation of aqueous suspensions of ultrafine particles," Colloids and Surfaces, 104(A) (1995), 101-109.

15. A.S. Koparal, and U.B. Ogutveren, "Removal of nitrate from water by electroreduction and electrocoagulation," Journal of Hazardous Materials, 89(B) (2002), 83- 94.

16. X. Chen, G. Chen, and P.L. Yue, "Separation of Pollutants from Restaurant Wastewaters," Seperation and Purification Technology, 19 (2000), 65- 76.

17. M. Kashefialasl et al., "K. Treatment of dye solution containing colored index acid yellow 36 by electrocoagulation using iron electrodes," International Journal of Environnemental Science Technology, 2(2006), 365-371.

18. M. Kobya et al., "Treatment of levafix orange textile dye solution by electrocoagulation," Journal of Hazardous Materials, 132 (2006), 183-188.

19. Y. Tamaura, et al., "The Synthesis of Green Rust II(FeI"i-FeII2) and Its Spontaneous Transformation into Fe304," Bulletin of the Chemical Society of Japan, 9 (1984), 2411-2416.

20. T. Satapanajaru et al., "Green rust and iron oxide formation influences metolachlor dechlorination during zero valent iron treatment," Journal of Environmental Science and Technology, 37 (22) (2003), 5219-5227.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

REMOVAL OF ARSENIC BY USING GREEN RUST AND OTHER ELECTROCHEMICALLY GENERATED FLOC

Md. Sanoar Rahman1, Jewel Andrew Gomes2, Kevin Urbanczyk3, David L. Cocke4

'Department of Civil Engineering, Lamar University, Beaumont, TX 77710 2Dan F. Smith Department of Chemical Engineering, Lamar University, Beaumont, TX

77710 'Department of Geology, Sur Ross State University, Alpine, TX

4Gill Chair of Chemical Engineering, Lamar University, Beaumont, TX 77710

Keywords: Electrocoagulation, copper, iron, aluminum, arsenic

Abstract

Arsenic contamination from groundwater is a worldwide problem. It is considered as one of the toxic materials being controlled by Environmental Protection Agencies in several developed and developing countries. The most common form of inorganic arsenic in natural water system is arsenate and arsenite. Arsenite (As(III)) is more toxic than Arsenate (As(V)) and it is more abundant in anaerobic condition such as in groundwater. Removal of arsenic from synthetic wastewater was carried out by using Green Rust floe which belongs to the layered double hydroxide (LDH) family and it was produced in-situ during electrocoagulation procedure. Other electrochemically generated floes were also used for the removal of arsenic from synthetic wastewater. The maximum arsenic removal efficiency was recorded as 99.9%. The sludge produced during electrocoagulation was assessed for semi-conductive property with diffuse reflectance UV-Vis spectroscopy. Determination of the energy gap (band gap) of the EC sludge provides a value of 1.43 eV that affirms the potentiality of EC byproducts as semiconductor.

Introduction

Arsenic is a toxic element which is present in natural water usually in trace concentrations. It is a major threat for the life of humans and useful microorganisms as concentrations become elevated. The national and international drinking water standards set by the World Health Organization, (WHO), is 10 ugL"' [5,7]. Arsenic concentration in soils and water can become elevated due to several reasons, like use of arsenical pesticides, mineral dissolution, disposal of fly ash, mine drainage, and geothermal discharge [1]. It is present in natural waters in both inorganic and organic forms. The most common form of inorganic arsenic in surface water is arsenate (As(V)), while under highly reducing conditions, such as anaerobic ground waters, arsenic may form reduced arsenite (As(III)) forms [2]. As (III) is more toxic than As (V). The concentration of arsenic species is mainly dependent on redox potentials [3] and pH [4]. Many people in the world are exposed to excessive amount of arsenic in drinking water [5,6]. The most serious problems being encountered in many regions of the world such as Argentina, Bangladesh, Chile, India, Mexico, Mongolia, Myanmar, Nepal, New Zealand, Thailand,

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Taiwan, Turkey, and Vietnam when arsenic concentration in ground and surface water exceeds national and international drinking water standards of 10 ugL"1 set by the WHO [5,7]. Several processes are available to treat wastewater, such as coagulation, membrane filtration, reverse osmosis, electrooxidation, electrofenton and finally electrocoagulation (EC). Among all the existing treatment technologies, EC received significant focus because of its cost effectiveness, ease for establishment and fewer amounts of byproducts. Electrocoagulation is a simple and efficient method where the flocculating agent is generated by electro-oxidation of a sacrificial anode. In this process, the treatment is performed without adding any chemical coagulant or flocculant, thus reducing the amount of sludge which must be disposed [8].

Experimental

EC was run both in flow-through EC apparatus (FTEA, manufacturer: Kaselco) and in a beaker-size reactor. The FTEA essentially consists of a flow-through cell, the electrode assembly, the feed pump and the DC power supply unit. The basic procedure was to pass the sample solution through the reactor within the vicinity of the electrodes inside the reactor or keep the solution in the vicinity of the electrodes while current applied. The volume of the reactor was 450 mL for the FTEA system. The beaker size EC was carried out in a 250 ml beaker with magnetic stirrer using vertically positioned electrodes dipped in the wastewater. The current and voltage during the EC process were checked using Cen-Tech multimeters. The pH of the solutions before and after EC was measured by an Oakton pH meter. EC was run for a various periods of time according to the experiment. After each run, the EC-mixture was filtered and the precipitate was dried. The filtrate was used for determining the amount of residual arsenic.

(T) Batch Scale Experiments Three combinations of sacrificial electrodes were used for the batch scale experiments: Cu-Cu, Cu-Fe and Cu-Al. Batch electrocoagulation experiments were carried out in a beaker with 150 ml solution using vertically positioned electrodes dipped in the solution. The dimensions of the anode and cathode were 3.2 cm x 1.93 cm, and the electrodes were submerged 2.5 cm under the sample solution. Both electrodes had a submerged surface area of 10.52 cm2. The electrode spacing was 2 cm. Electrodes were connected to a DC power supply (Kaselco power supply). Electrocoagulation experiments were performed at various applied current level. A digital voltmeter was used to measure the cell voltage. It is shown in Figure 1.

Electrocoagulation experiments were run for different residence time ranging from 5 minutes to 60 minutes. The polarity of the electrodes was changed to improve the performance of EC after half of residence time for each experiment. Prior to each experiment, the electrodes were abraded with sand-paper to remove scale formed in previous runs then treated with a solution of 10% HNO3 in order to remove any effect due to the different prehistory of the electrodes, washed with distilled water, dried and then, the anode was weighted. Experiments were conducted with lppm Arsenic solution. Samples were drawn from the solution after completion of each experiments (residence time 5, 10, 20, 30 and 60 minutes), filtered and analyzed for arsenic.

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Figure 1. Schematic of the Batch Reactor.

(ii) Flow-through EC Apparatus (FTEA)added a blank line

A flow-through EC apparatus (FTEA) was supplied by Kaselco EC Treatment System, Shiner, Texas, USA. The FTEA essentially consists of a flow-through cell, the electrode assembly, the feed pump and the DC power supply unit. A schematic diagram of the FTEA is shown in Figure 2.

The cell contained five parallel electrodes (11.0 cm x 11.4 cm) placed 6.0mm apart, which forms four parallel cells. Before use the plates were cleaned manually by abrading with sand paper. The internal volume of the cell is approximately 450 mL. A variable transformer was used to control the current and the applied potential [9]. In a typical EC experiment using the FTEA apparatus the sample was pumped through at a predetermined flow rate of 505.05 mL/min giving a residence time of 3.56 minutes. The solution in an amount of 2.5 liter was treated in each batch.

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Figure 2. Diagram of the Kaselco Bench Reactor.

Results and discussion

Arsenic Removaladded a blank line

Arsenic removal was affected by residence time of the experiments and the electrodes materials used. There was a sharp increase in the removal efficiency of arsenic from the sample in the lower levels of residence time. The removal efficiency was in the range of 56.2% to 66.2% when the residence time was 5 minutes for all electrode combination. This efficiency increased to a range of 97.5% to 99.9% when the residence time was increased from 5 minutes to 30 minutes. The removal of Arsenic was rapid in the first 5 to 10 minutes of residence time. The removal continued to increase up to 30 minutes of residence and then began to level out. This can be explained by two reasons. First of all there are more abundant arsenic ions at the beginning of the EC experiments and therefore the rate of reaction and adsorption is high whereas there is a slower reaction rate at the end of the process as a result of reduced arsenic concentration. Secondly, the sludge that produced during the EC process covers some surface area of the electrodes which decreases the efficiency of arsenic removal. The permissible limit of arsenic concentration in the drinking water was attained with a combination of Cu-Cu and Cu-Fe electrodes with a residence time of 30 minutes. There was a little increase in removal efficiency after a residence time of 30 minutes. Considering the cost of operation of the treatment process a residence time of 30 minutes was optimum for all electrodes combination.

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Figure 3. Effect of Residence Time on the Removal Efficiency of Arsenic.

Reaction Kineticsadded a blank line

The goal of any particular kinetics experiments is to measure the concentration of a particular chemical species over time and in this way the rate law can be determined. However, it is very difficult to measure the concentration of a particular species at a particular time because of the methods or technique used to measure the concentration do not work instantaneously. An idea about the order of reaction can be determined if the concentration of the species is plotted against time even the measurements are not most accurate. A graph of the the arsenic concentration was plotted versus time (Figure 4) and it showed the same trend as first order reaction. The graph of ln[As cone] was plotted against time and as for other first order reaction the plot was a straight line (Figure 5) with a negative slope.

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Figure 4. Arsenic Concentrations versus Time in minutes.

Figure 5. ln[As conc] versus Time (min.) plot.

For first order reaction:

l n r r - = ~kt (1)

Where, A = concentration at time t Aç> = concentration at time 0 (initial concentration) k = rate constant. From Figure 5, k was determined as 0.22 min"1.

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Semi-conductive property of the produced sludgeadded a blank line

Energy gap (Eg) is an important feature of semiconductor. The sludge produced during electrocoagulation was assessed for semi conductive property with diffuse reflectance UV-Vis spectroscopy. The intersection between the linear fit and the photon energy axis (eV) gives the value of Energy gap. Determination of the energy gap (band gap) of the EC sludge produces a value of 1.43 eV. Most common semiconductors has energy gap value between 1.11 eV to 2.86eV ( Si (1.1 leV); Se (1.74 eV); SiC (2.86eV)), that states the potentiality of EC byproducts as semiconductor.

Figure 6: Determination of the band gap of the EC sludge giving a value of 1.43 eV.

Conclusion

The use of dissimilar electrodes and higher conductive electrodes, such as copper, provides alternative method for the removal of arsenic by electrocoagulation from wastewater. The removal efficiency was found to be affected by generated floe from the electrodes materials used and also residence time. The maximum arsenic removal efficiency was determined as 99.9% from the initial arsenic concentration of 1 ppm. The arsenic removal profile with time follows first order kinetics with a rate constant value of 0.22 min"1. The EC-sludge can be reused as valuable byproducts in semi-conductor industries.

Acknowledgements

We greatly acknowledge the financial support from USDA (2009-38899-20017), and the instrumental support from Lamar Materials Instrumentation Center.

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References

1. M.Y.A. Mollah, R. Schennach, J.R. Parga, and D.L. Cocke, "Electrocoagulation (EC)-science and applications,"/ Hazard. Mater., 84 (2001), 29^tl.

2. M.J. Kim, J. Nriagu, and S. Haack, "Arsenic species and chemistry in groundwater of southeast Michigan", Environ. Pollut., 120 (2002), 379-390.

3. J. G. Hering, P.Y. Chen, J.A. Wilkie, M. Elimelech, and S. Liang, "Arsenic removal by ferric chloride," J. Am. Water Works Assoc, 88 (1996), 155-167.

4. Masscheleyn et al., "Effect of redox potential and pH on arsenic speciation and solubility in a contaminated soil," Environ.Sci. Technol., 25 (1991), 1414-1419.

5. P. Ravenscroft, H. Brammer, and K. Richards, Arsenic Pollution: A Global Synthesis, RGS IBG Book Series (United Kingdom: A John Wiley & Sons Publication Ltd., 2009).

6. P. L. Smedley, and D.G. Kinniburgh. "A Review of the Source, Behavior and Distribution of Arsenic in Natural Waters," Applied Geochemistry, 17 (2002), 517.

7. WHO (World Health Organization), Guidelines for Drinking Water Quality 1 (2006), 306. 8. Cenkin et al., "Electrochemical treatment of industrial wastewater," Eff. Water Treat. J., 25 (7) (1985), 243-249.

9. M. Y.A. Mollah, P. Morkovsky, J.A.G. Gomes, M.Kesmez, J. Parga, and D.L. Cocke, "Fundamentals, Present and Future Perspectives of Electrocoagulation," Journal of Hazardous Materials, Bl 14 (2004), 199-210.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

FORMATION OF LAYERED DOUBLE HYDROXIDES IN SELF-PURIFICATION OF POLYNARY METAL

ELECTROPLATING WASTEWATERS FOR EFFECTIVE REMOVAL OF ANIONIC DYE

Ji Zhi Zhou, Guangren Qian, Chong Liu, Yueying Wu, Xiuxiu Ruan, Yunfeng Xu, Jianyong Liu

School of Environmental and Chemical Engineering, Shanghai University; 333 Nanchen Rd.; Shanghai 200444, PR China

Keywords: Layered double hydroxide(LDH); Heavy metal; Anionic dye; Adsorption

Abstract

The synchronous formation of layered double hydroxide (LDH) precipitates was performed for the self purification of wastewaters containing Ni2+, Zn2+, and Cr3+. With ~ 99% of metal ions removed, a pure Ni-Zn-Cr-bearing LDH was formed by accelerated carbonation process. The obtained LDH were characterized by XRD, FT-IR, SEM and BET surface area techniques. Moreover, the removal of a dye, Acid Scarlet GR (GR), from aqueous solution on the LDH was studied under varying conditions of pH, adsorbent dose and contact time. It is observed that Ni-Zn-Cr-LDH could effectively remove GR from aqueous solutions and the removal capacity was increased with rising temperature. The low concentration of heavy metal in the final solution after dye removal indicates there was little heavy metal dissolved from as-obtained LDH. These results suggest a strategy to treat heavy metal wastewater efficiently and propose Ni-Zn-Cr-LDHs as an environmental friendly adsorbent for anionic dye removal.

Introduction

Heavy metals-bearing wastewater is a major concern for the environment as heavy metal was toxic and hazardous to human health and ecosystem. The chemical precipitation using calcium (Ca) sources such as lime is widely used to remove heavy metals in wastewater [1], in which heavy metals usually form the insoluble hydroxides or calcium-heavy metal mixed hydroxides as lime dissolution [2]. However, this lime precipitation method results in final pH usually ranging from 9 to 11 [1], higher than that in the discharge limit [3]. Moreover, the most difficult issue emerges from the large amount of solid residue which needs to be disposed.

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Accelerated carbonation technology as an advanced treatment with calcium sources has been developed to treat heavy metals from waste streams [4-7]. Most of the heavy metals can be precipitated by pumping carbon dioxide to form insoluble heavy metal carbonates together with CaCC>3. The benefits of the carbonation method are the constant pH of effluent in a lower range (pH = 8-10) after the neutralization effect of CO2 [8] as well as the heavy metal carbonate with high solubility product. Despite the significant promotion of heavy metal removal, it is still a challenge for carbonation technology to treat the formed solid waste residue that is composed of a large amount of CaCÛ3 [9, 10]. Consequently, the formation of this by-product CaCC"3 does not remarkably decrease the amount of solid residue after treatment [5,11].

Besides the hydroxide and carbonate, heavy metals can also form the heavy metals-layered double hydroxide (HM-LDH) which could be usually in natural soil and hydration product from the heavy metal solidification/stabilization by cement concrete [12, 13]. A common formula of LDH is [M2+,.xM

3+x(OH)2]

x+(A„.)x/n • mH20, where M2+ is bivalent metal, M3+ tri valent metal, An" anion, and x the molar ratio of trivalent metal to total metal. Generally, most of M2+ and M3+ can be incorporated in LDH materials, including Co2+, Fe2+, Ni2+, Cd2+, and Zn2+ as M2+ while Fe3+ and Cr3+ as M3+

[14, 15]. Besides the binary-LDH, it is feasible to obtain polynary-LDH where more than two metal cations are introduced [16]. Moreover, the molar ratio of M2+/ M3+ can be changed from 2.0 to 4.0 (corresponding to x = 0.20-0.33) [17]. On the other hand, LDHs containing heavy metal are promising materials for many practical applications in catalysis, adsorption, pharmaceutics, photochemistry, electrochemistry, and other areas [18]. Accordingly, it is possible to induce the co-precipitation of heavy metals with distinct valency to form various polynary HM-LDHs (PHM-LDHs) via the modulation of M2+/ M3+ with variable molar ratios without Ca addition. Furthermore, the produced LDH could be an attractive material as its high adsorption capacity of anions, which is an advantage for removing many pollutants [19].

Herein, we report an accelerated carbonation without calcium addition to directly purify a heavy metal containing solution with diverse ratios of Zn2+, Ni2+, and Cr3"1". We found that most of heavy metals were removed via the formation of pure PHM-LDH. Potential application of the products was also performed via the adsorption of acid brilliant scarlet GR dye. The strategy developed in this work is expected as a new promising prospect for environmental friendly removal and recovery of heavy metals in wastewater.

Experimental Section

Heavy metals removal and LDH synthesis The simulation experiments for the rapid accessibility were carried out with various

initial concentrations of heavy metals (Table I), reaction temperature, pH value and reaction time. Typically, 50 ml of the wastewater was added into 50 ml distilled water under stirring at 25 °C for an hour. Synchronously, LDH containing heavy metal ions

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was prepared by a co-precipitation method by bubbling a stream of C0 2 into the wastewater at a constant rate of 200 ml/min and adjusting the pH of reaction solution to 8.0 with 0.28 M NaOH. After reaction, the suspension was transferred into a conical flask and aged at 60 °C for 18 hr with sequentially shaking. The solid product was filtered and washed by distilled water. After dried at 60 °C for 24 hr, solid was ground for further characterization. LDH products prepared are denoted as LDH332, LDH222, et al., respectively, as corresponding to the heavy metal ratio in the simulated solution.

Characterization X-ray diffraction (XRD) patterns were collected by Rigaku DLMAX 2200 with Cu

Ka (X = 0.154 nm) radiation at 40kV. The lattice parameters for LDH product were calculated by MDI Jade software 5.0 with ICDD database 2003. The inductively coupled plasma atomic emission spectrometry (ICP-AES, LEEMAN Prodigy) was used to determine the concentration of heavy metals in solution with 1 mol/L nitric acid as digesting reagent. The concentration of GR in aqueous solution was determined by a spectrophotometer (UV-2450, SHIMADZU) at 510 nm [20].

Adsorption ofGR Dye GR (C22Hi4N4Na207S2, 99.9%) was purchased from Anhui Fengyang dye

chemical Co. Ltd, China. Isothermal adsorption experiment was performed by adding 0.1 g as-obtained LDH into 50 ml GR solution with different concentrations and stirring for 24 hr before taking a 2 ml of liquid sample for concentration analysis. Langmuir equation was used to fit the experimental equilibrium isotherm data [21].

Kinetic experiment was conducted using the same dose of the LDH into 50 ml of solution containing 100 mg/L of GR. The solution was separated in the same way in the case of isotherm experiment. The concentration of residual dye GR was determined at selective time. All adsorption experiments were carried out with sequential shaking in water bath.

Results and Discussion

Attractiveness of Self-Purification For most of LDH synthesis, the ratio of trivalent cation to total cations (x) varied

from 0.20 to 0.33. Some researchers tried to extend the x value scale in the synthesis of LDH. For some binary LDHs such as ZnAI-LDH and ZnCr-LDH, x value has been developed up to 0.46 [17].

In the present work, we tried to treat more complicated wastewaters to form LDH with the x ratio ranging from 0.2 to 0.4. To estimate the effect of varying x on the formation of polynary heavy metal LDHs, simulation tests were designed using Ni, Zn, and Cr containing solutions with various initial heavy metal ratios. As shown in Figure 1, all synthesized samples show the characteristic XRD patterns of LDH structure that

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is indexed as a series of reflections (003),(006),(012), and (110). Such evidences indicate a pure LDH phase formation with various x values.

Figure 1. XRD patterns for as-obtained products from the purification of simulated wastewater with various initial heavy metal ratios

Table I. Lattice parameters of LDHs obtained from simulated wastewater with various molar ratios of heavy metals

X

value

0.20

0.25

0.33

0.40

Ni:Zn:Cr (theoretical)

2:2:1 1:3:1 3:1:1 3:3:2 4:2:2 2:4:2 5:1:2 1:5:2 2:2:2 1:3:2 3:1:2 2:1:2 1:2:2

Ni:Zn:Cr (experimental)

2.0 1.1 3.2 3.1 4.2 2.1 5.2 1.1 2.2 1.1 3.3 2.1 1.1

2.3: 1.0 3.6: 1.0 1.1: 1.0 3.3: 2.0 2.3: 2.0 4.2: 2.0 1.0 :2.0 5.1:2.0 2.5 :2.0 3.8: 2.0 1.1:2.0 1.1:2.0 2.6: 2.0

doo3 (A)

7.62 7.60 7.66 7.66 7.72 7.66 7.92 7.61 7.65 7.64 7.88 7.84 7.78

a (nm)

0.309 0.310 0.308 0.308 0.306 0.309 0.307 0.310 0.309 0.311 0.308 0.308 0.308

c (nm)

2.28 2.30 2.29 2.28 2.32 2.32 2.25 2.30 2.31 2.30 2.35 2.36 2.31

Reduction percentage

(%) Ni

99.6 99.4 99.5 99.4 99.4 99.4 99.5 98.9 99.4 99.5 99.3 99.1 99.7

Zn 99.6 99.8 99.7 99.6 99.7 99.7 99.7 99.8 99.8 99.8 99.6 99.8 99.6

Cr 99.8 99.8 99.8 99.9 99.9 99.9 99.9 99.9 99.7 98.9 99.6 99.6 99.7

Moreover, Table I lists the initial molar ratios of heavy metals in solutions and the lattice parameters for the resultant products. The calculated interlayer space dow was varied regularly with the changing of molar ratios of heavy metals (Table I). For instance, doo3 increased from 7.60 to 7.88 with the value of x increasing from 0.2 to 0.4. On the other hand, molar ratio of bivalent metals also has an effect on the changing of oW For example, the doo3 reduced from 7.92 to 7.61 when the ratio of Ni/Zn was decreased from 5/1 to 1/5. Thus, we expect that the LDH structure can be controlled by changing heavy metal ratio.

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After LDH formation, the heavy metal composition of product was close to that of theoretical LDH in most cases (Table I). As listed in Table I, it is contributed to about 99% of the removal percentage of each heavy metal. As a result, we expect that the self-purification process of heavy metal-containing wastewater concurrent with the formation of LDH is easily performed in wide conditions.

Effect of operational conditions Typically, LDH preparation demands the high alkalinity in solution [14]. Boclair and

Brateman observed the precipitation of LDH containing Cr occurred directly by a one-step process under low pH without the formation of a precursor M3+ hydroxide phase, in contrast to the related systems where M3+=A1 or Fe [15]. In our work, PHM-LDH was obtained at pH=6.0 (Figure 2A), which might be attributed to the formation of Cr-OH-M(II) bridge from deprotonated Cr monomers and M(II) in acidic condition [22]. The XRD pattern of the PHM-LDH obtained at pH=6.0 is same as that of PHM-LDH obtained at a higher pH (Figure 2A). It represents that the current method can be used to prepare PHM-LDH from heavy metals-containing industrial wastewaters at a low pH without any additional calcium. The simulation experiments were also carried out under various temperatures with different reaction times while initial molar ratio of heavy metals in solution was kept at Ni:Zn:Cr=3:3:2. As shown in Figure 2, the pure LDH structure is observed from the XRD patterns of the products. With an increasing temperature from 25 °C to 80 °C, the peaks of XRD patterns become sharper (Figure 2B), indicating that higher temperature leads to better crystallinity. XRD patterns (Figure 2C) of the PHM-LDHs obtained after 1 hr or 72 hr reaction show very similar peaks, indicating a rapid self-purification occurred in the formation of PHM-LDH.

Figure 2. XRD patterns of LDHs prepared from the simulation wastewaters with an initial molar ratio of heavy metals Ni:Zn:Cr = 3:3:2 under various synthesis conditions: (A) temperature, (B) pH, and (C) reaction time.

The two key factors (pH and time) can control not only the formation of ZnNiCr-LDH but also the removal efficiency of heavy metals in the self purification process. The remarkable reduction of heavy metals in wastewater occurred within initial one hour. The concentration of Zn decreased from 735 mg/L to 2.93 mg/L, Ni

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from 660 mg/L to 2.10 mg/L and Cr from 389 mg/L to 1.25 mg/L. After 10 hour reaction, the concentration of Zn was 0.118 mg/L, Ni 0.140 mg/L and Cr 0.116 mg/L. And these concentrations still remained the same even after 24 h reaction.

When pH was kept within 8.0 to 10.0, the remaining concentrations of heavy metals were lower after the carbonation. In such pH range, the Zn concentration was changed from 0.423 mg/L to 0.199 mg/L, Ni from 0.400 mg/L to 0.359 mg/L, and Cr from 0.267 mg/L to 0.351 mg/L. With the decreasing of pH down to 5.0, Zn and Ni were largely released up to 91.9 mg/L and 167 mg/L respectively. The same situation was observed when the pH was higher than 10.0. The release of heavy metals is due to the dissolution of LDH at lower pHs and the soluble fraction Zn(OH)4

2" at higher pH. Thus, pH from 8.0 to 10.0 is optimum for preparing the LDH products, which leaves the least amount of heavy metals in solution.

Adsorption ofGR onto LDH product. As a promising material for pollutants adsorption, LDH has been applied for the

removal of organic dye [23]. In this work, an adsorption of azoic dye GR onto LDH 332 was evaluated. The adsorption kinetics is depicted in Figure 3A, indicating that a multiplex model (solid line) gave a good fit to the experimental data [24], In the adsorption process, 94.4% of GR was reduced with its initial concentration of 100 mg/L within 24 h. It indicates the feasibility of PHM-LDH product in removing GR. Figure 3B illustrates the adsorption isotherms of GR on LDH 332 at two different temperatures (25 and 30 °C). The equilibrium adsorption isotherms at 25 and 30 °C were well fitted by Langmuir adsorption model (R2>0.999, Table II). Increasing temperature could improve the adsorption capacity of LDH. At 30 °C, the adsorption amount of dye reached to 49.7 mg/g.

The determination of heavy metals in the final solution after dye adsorption showed the concentrations of Ni and Cr remained at about 0.04 mg/L and 0.15 mg/L during 48 hr adsorption respectively. The concentration of Zn was about 0.16-0.32 mg/L. It indicates that less heavy metal was leaked from LDH in the adsorption process with pH varying between 6.0 and 8.0. The stability of structure and less heavy metal releasing demonstrate a promising performance and friendly environmental application of PHM-LDH for adsorption of organic dye.

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Figure 3 Adsorption of Acid Brilliant Scarlet GR on the LDH 332: (A) adsorption kinetics at 25°C; (B) Langmuir adsorption isotherm at 25 and 30 °C, respectively.

Table II Thermodynamic parameters for the adsorption of GR on PHM-LDH

Temperature

(°C)

25 30

KL

(L/mg)

1.963 1.491

Qm (mg/g) 44.41 50.38

R2

0.99995 0.9997

Acknowledgments

This project isfinancially supported by National Nature Science Foundation of China No. 20907029, 20877053 and No. 20677037, Shanghai Leading Academic Discipline Project No. S30109. We appreciate Instrumental Analysis & Research Center of Shanghai University for the XRD and SEM/EDX determination.

References

1. T.A. Kurniawan et al., "Physico-chemical treatment techniques for wastewater laden with heavy metals," Chem EngJ, 118(1-2) (2006), 83-98. 2. X.M. Lin et al., "Heavy metals in wastewater: The effect of electrolyte composition on the precipitation of cadmium(II) using lime and magnesia," Water Air SoilPollut, 165 (1-4) (2005), 131-152. 3. Ministry of Environmental Protection of PRC, Discharge standard of pollutants for municipal wastewater treatment plant GB 18918-2002(ßeijmg, 2002). 4. M.F. Bertos et al., "A review of accelerated carbonation technology in the treatment of cement-based materials and sequestration of CO2," J Hazard Mater, 112 (3) (2004), 193-205. 5. J.C. Walton et al., "Role of carbonation in transient leaching of cementitious wasteforms," Environ Sei Technol, 31 (8) (1997), 2345-2349. 6. M.F. Bertos et al, "Investigation of accelerated carbonation for the stabilisation of MSW incinerator ashes and the sequestration of CO2," Green Chem, 6 (8) (2004), 428-436. 7. G. Montes-Hernandez et al., "Removal of oxyanions from synthetic wastewater via carbonation process of calcium hydroxide: Applied and fundamental aspects," J Hazard Mater, 166 (2-3) (2009) 788-795. 8. X. Li et al , "Accelerated carbonation of municipal solid waste incineration fly ashes," Waste Manage, 27 (9) (2007), 1200-1206. 9. Q.Y. Chen et al., "Precipitation of heavy metals from wastewater using simulated flue gas: Sequent additions of fly ash, lime and carbon dioxide," Water Res, 43 (10) (2009), 2605-2614.

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10. G. Montes-Hernandez et al., "Mineral sequestration of C0 2 by aqueous carbonation of coal combustion fly-ash,". J Hazard Mater, 161 (2-3) (2009), 1347-1354.

11. J.-g. Jiang et al., "Continuous CO2 capture and MSWI fly ash stabilization, utilizing novel dynamic equipment," Environ. Pollut, 157 (11) (2009), 2933-2938. 12. I. Moulin et al., "Retention of zinc and chromium ions by different phases of hydrated calcium aluminate: A solid-state Al-27 NMR study," J Phys Chem B, 104 (39) (2000), 9230-9238. 13. M. Vespa et al., "Spectroscopic investigation of Ni speciation in hardened cement paste," Environ Sei Technol, 40 (7) (2006), 2275-2282. 14. Z.P. Xu et al, "Dispersion and size control of layered double hydroxide nanoparticles in aqueous solutions," J Phys Chem B, 110 (34) (2006), 16923-16929. 15. J.W. Boclair et al. "Layered double hydroxide stability. 2. Formation of Cr(III)-containing layered double hydroxides directly from solution," Chem Mater, 11 (2) (1999), 303-307.

16. Y. Guo et al., "Synthesis and characterization of Cd-Cr and Zn-Cd-Cr layered double hydroxides intercalated with dodecyl sulfate," J Solid State Chem, 178 (6) (2005), 1830-1836. 17. P.S. Braterman, Z.P. Xu, F. Yarberry. "Layered Double Hydroxides(LDHs)," Handbook of Layered Materials, ed. S.M. Auerbach, K.A. Carrado, P. K. Dutta (New York, NY: Marcel Dekker, 2004), 377. 18. F. Zhang et al., "Layered Double Hydroxides as Catalytic Materials: Recent Development," Catal Surv Asia, 12 (4) (2008), 253-265. 19. K.H. Goh, T.T. Lim, Z. Dong, "Application of layered double hydroxides for removal of oxyanions: A review," Water Res, 42 (6-7) (2008), 1343-1368. 20. X. Wang, N. Zhu, B.Yin, "Preparation of sludge-based activated carbon and its application in dye wastewater treatment," J Hazard Mater, 153(1-2) (2008), 22-27. 21. S. Mandai et al, "Azoic Dye Hosted in Layered Double Hydroxide: Physicochemical Characterization of the Intercalated Materials," Langmuir, 25 (18) (2009), 10980-10986.

22. H. Roussel et al., "Study of the formation of the layered double hydroxide [Zn-Cr-Cl]," Chem Matel, 13 (2) (2001), 329-337. 23. F.B.D. Saiah, B.L. Su, N. Bettahar, "Nickel-iron layered double hydroxide (LDH): Textural properties upon hydrothermal treatments and application on dye Sorption," J Hazard Mater, 165 (1-3) (2009), 206-217. 24. L. Lv et al., "Treatment of high fluoride concentration water by MgAl-CCb layered double hydroxides: Kinetic and equilibrium studies," Water Res, 41 (7) (2007), 1534-1542.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

CHARACTERIZATION AND CHEMICAL MODIFICATION OF ELECTROCHEMICALLY PRODUCED LAYER DOUBLE HYDROXIDES

AS NANOMATERIALS

Md. Kamrul Islam1, Jewel Andrew Gomes2, Paul Bemazzani1

'Department of Chemistry and Biochemistry 2Dan F. Smith Department of Chemical Engineering

Lamar University, Beaumont, TX 77710

Keywords: Layered Double Hydroxides, electrocoagulation floe, Nanomaterials, nanowire

Abstract

Electrocoagulation (EC) is one of the most effective electrochemical techniques for the removal of trace amount of metal contaminants from waste water. The intermediate step of this process uses a sacrificial iron electrode, generating Layered Double Hydroxides (LDHs) known as green rust. When trace copper is removed using EC, the produced EC-floc contains mostly iron with a smaller proportion of copper incorporated in it. Instead of disposing EC-floc as waste, further treatment produces important nanomaterials. The EC floe treated chemically convert a portion into nanowires of iron and copper. The proportion of copper incorporated into iron floe was varied to observe the effect of change of the ratio to copper in the structure, size, shape of nanowires. Copper was recaptured both chemically and electrochemically, and converted to nano rods. The nanomaterials were characterized using SEM and XRD.

Introduction

Electrocoagulation is one of the most successful and cost effective traditional electrochemical techniques over the past several decades for waste water treatment. This technique is most effective for contaminant metallic ion removal from waste water. Copper is a commonly used metal used for commercial and household usage. Overuse of this material leads to a significant amount of it in water. Copper was removed successfully (about 99.99%) by Electrocoagulation processes using iron electrodes [1]. Floe that is generated during the EC process is responsible for the removal of copper from the copper contaminated waste water. The method involves the precipitation of ions of interest by the addition of counter ions that destabilize the material to be extracted. In this study, iron was used to destabilize the copper solution and cause precipitation. The precipitate is a mixture of iron and copper. When the trace copper is removed by EC, the produced EC floe contains mostly iron and lesser proportion of copper incorporated in it. Instead of treating the EC floe as waste, it was further treated to produce important nanomaterials, like nanowires.

The advancement of nanotechnology continues to impact our daily lives, especially innovations in the electronics industry. Nanowires are expected to play a vital role in future integrated circuit applications. Iron and copper are excellent candidates for this role due to their high electrical conductance, abundance and low cost. The utilization of electrocogulation waste, i.e. EC floe for the conversion or synthesis of nanowires can be a useful technique for waste management

445

processes. Both chemical and electrochemical processes were used for the conversion of the EC floe into nanowires. As the proportion of copper decreases in EC floe, copper chloride (CuCk) solutions were used to compare chemical and electrochemical production of Cu nanowires.

Layered Double Hydroxide (Hydrotalcite) Layered double hydroxide (LDH) is one of the most promising materials, widely used as catalyst, catalyst support, adsorbent with large anion adsorption capacity. LDH can be considered as anti-types of 2:1 (three-layer) clay minerals [2]. The general chemical formula is [M"i-IM

IIII(OH)2]jr+[A""r/„]I"^H20, where M11 is a divalent cation (Mg2+Zn2+, Mn2+, Mo2+, Ni2+,

Cd2+), M1" is a trivalent cation (Al3+, Fe3+, Cr3+, Ga3+), A"" is a anion, such as carbonate, chloride, or sulfate. Trivalent cations (e.g., Al3+) generate positive charges on the layers, which are compensated by anions located in the interlayer region. Water molecules are also found in the interlayer space [3]. In a brucite layer of hydrotalcite, each Mg2+ ion is octahedrally surrounded by six OH- ions, with the different octahedra share edges forming an infinite two-dimensional layer. The brucite crystals are then formed when the 2-D layers are stacked one on the top of another via 0-H--0 hydrogen bonding between the adjacent layers. Partial replacement of Mg2+

ions by Al3+ gives the brucite-like layers a positive charge, which in hydrotalcite itself is balanced by carbonate anions, located in the interlayer region (gallery) between the two brucite-like layers. This gallery also contains water molecules, hydrogen bonded to layer OH" ions and/or to the interlayer anions. The electrostatic interactions and hydrogen bonds between the brucite-like layers and the gallery species hold the layers together, forming the three-dimensional structure [4]. Hydrotalcite (HT) compounds can be viewed as a stacking of positive sheets, a brucite-like structure, where metallic cations are located in the layer, while anions and water are found in the interlayer space. Possibile introduction of active metals, such as Pd, Cu, and Ni, into the layers can produce high dispersion, may increase the stability and activity of the catalyst. Due to the special structural arrangement of the interlayers or galleries, it may be hypothesized that toxic anions (such as nitrate) in contaminated water can be concentrated in the interlayer space and be removed from the aqueous phase.

Experimental

Chemical synthesis of nanowire: (A). In a typical preparation procedure, copper chloride (0.17g, 12.5mM) and glucose (0.391 g, ~2mmol) were massed and dissolved in 80 mL of distilled water in an Erlenmeyer flask. Hexadecalamine (1.44 g) was slowly added to the water solution followed by vigorous mixing for 6 h with a magnetic stirrer until a light blue emulsion was obtained. The emulsion was then placed in a Teflon-lined pressure reactor of 200 mL capacity. The pressure reactor was heated for 2 h at 393 K under autogenous pressure and allowed to cool to room temperature. The resulting reddish brown solution was centrifuged (2000 rpm) and washed with deionized water, n-hexane and ethanol sequentially. The process was repeated several times to remove the excess surfactant, and reddish fluffy solid. The product was kept under n-hexane to avoid oxidation of the copper nanowires [5],

(B). For the preparation procedure of nanowire from the copper containing iron EC floe, EC floe (0.34g) and glucose (0.391 g, ~2mmol) were dissolved in 80 mL of distilled water in an

446

Erlenmeyer flask. Hexadecaylamine (1.44 g) was slowly added to the solution followed by vigorous mixing for 48 h with a magnetic stirrer until a blackish shiny colored emulsion was obtained. The emulsion was then placed in a Teflon-lined pressure reactor of 200 mL capacity. The pressure reactor was heated for 5 h at 393 K under autogenous pressure and cooled to room temperature. The resulting brownish black solution was centrifiiged (2000 rpm) and then washed with deionized water, n-hexane and ethanol sequentially. The process was repeated several times to remove the excess surfactant, and brownish solid was obtained. The product was kept under n-hexane to avoid oxidation of the nanowires.

Electrochemical synthesis of nanowires (C). The back surfaces of the anodic aluminum oxide templates (AAO) were coated with a thick layer of gold using a sputtering spray. This AAO template was used as the working electrode during the electrodeposition of copper on it, using a three electrode potentiostat system (Voltalab) composed of a platinum wire counter electrode and Ag/AgCl reference electrode. Deposition was carried out at constant potential (-0.45 mV) for 1-1.5 h from a solution of 0.2 M CuCb 2H20, and 0.1 M boric acid in AAO template. The AAO template was dissolved for 1 h in 3 M sodium-hydroxide solution and copper nanorods were obtained [6].

(D). Copper was chemically separated from the EC floe with qualitative separation techniques [4], This chemically separated copper solution was used for the electrodeposition of copper to form copper nanowire. Like the previous method, gold coated AAO was used as working electrode for the deposition of copper using the three electrode potentiostat system (Voltalab) with a platinum wire counter electrode and Ag/AgCl reference electrode. Deposition was carried out at constant potential (-0.45 mV) for 1-1.5 h on AAO template [7]. The AAO template was dissolved for 1 hin 3 M sodium-hydroxide solution and copper nanorods were obtained.

Results and Discussion

Characteristics of Nanowire The prepared nanowires were characterized using SEM and XRD. SEM images (Figures 1, 3, 5 and 7) show the formation of nanowires from the electrically produced LDHs from the electrocoagulation techniques while removing copper from water using iron electrodes. XRD patterns (Figures 2,4, 6, and 8) show presence of iron and copper species in the chemically and electrochemically conversion of LDHs.

447

Figure 1. SEM image (A) shows the formation of nanowires from the CuCb solution. Extraction of nanowires from the reaction mixture is challenging for the chemical synthesis process.

Figure 2. X-ray diffraction pattern of the copper nanowire from CuCb solution. Panel (A) shows the presence of Cu in the nanowires synthesized from CuCU solutions. The broad peak

indicates amorphous regions present in the nanowires.

448

Figure 3. SEM image (B) shows the formation of chemically obtained nanowires from the Cu containing iron floe. The size differences suggests the floe nanowires are not pure Cu.

Fig. 4: X-ray diffraction pattern of chemically obtained nanowires from EC floe. The reference patterns; red: maghemite, blue: copper oxide, light green: goethite.

449

Figure 5. SEM image (C) shows nanowires synthesized from the electro-deposition of Cu on AAO from CuCh solution. The nanowires are relatively uniformly shaped.

Fig. 6: X-ray diffraction pattern shows the presence of Cu and CuO for the nanowires that was deposited electrochemically from a CuCh solution.

450

Figure 7. SEM image (D) shows electrochemically obtained copper nanowire from the quantitatively separated Cu from iron floe. In both cases, these nanowires have lower aspect

ratios.

Figure 8. X-ray diffraction pattern (D) shows copper nanowires which was electro-chemically deposited from the qualitatively separated Cu solution of Cu containing iron floe.

451

Conclusion

The synthesis of metallic nanowires by both chemical and electrochemical methods was successfully carried out. This exploratory work showed that Electrocogulation floe can be modified or converted to valuable nanowires by chemical and electrochemical techniques. SEM demonstrates the formation of nanowires. The EC method yielded nanowires of significantly lower aspect ratio. XRD shows the presence of crystalline Cu and CuO in the electrochemically deposited nanowires. Results show that nanowires can be synthesized chemically and electrochemically from EC floe, which is otherwise construed as waste.

Future Work Additional experiments are being performed to optimize conditions for the formation of uniform nanowires of larger aspect ratios. The amount and nature of iron in the nanowires synthesized from the floe will be investigated.

Acknowledgment We express our gratefulness to Welch foundation (Department of Chemistry, Lamar University) for their generous support for this research. In addition, we are thankful to Mr. Dan Rutman of Lamar Materials Instrumentation Center for allowing us to use SEM-EDS instruments. We also thank Dr. T. Thuy Minn Nguyen, Mr. Manohor Palla, for some experiments during the whole session.

References

1. J. Gomes et al., "Utilization of Electrochemiccal Techniques for copper Removal, Speciation, and Analysis in Aqueous System", Electrochemical Society Transactions, 28 (18) (2010), 59-68.

2. G. Grünewald, K. Kaiser, and R. Jahn, "Hydrotalcite—A potentially significant sorbent of organic matter in calcareous alkaline soils," Geoderma, 147 (2008) 141-150.

3. D. Wan et al., "Role of the Mg/Al atomic ratio in hydrotalcite-supported Pd/Sn catalysts for nitrate adsorption and hydrogénation reduction," Journal of Colloid and Interface Science, 332 (2009), 151-157.

4. P.S. Braterman, Z.P. Xu, and F. Yarberry. In: S.M. Auerbach, K.A. Carrado and P.K. Dutta, Editors, Handbook of Layered Materials (Dekker, New York, 2004), 373-474.

5. P.M. Ajayan, et. al., "Low-Temperature Large-Scale Synthesis and Electrical Testing of Ultralong Copper Nanowires", Langmuir, 26 (21) (2010), 16496-16502.

6. P.M. Ajayan, et. al., "Synthesis of Catalytic Porous Metallic Nanorods by Galvanic Exchange Reaction", J. Phys. Chem. C, 114 (2010), 389-393.

7. X.M. Liu, and Y.C. Zhou, "Electrochemical deposition and characterization of Cu20 nanowires," Appl. Phys. A, 81 (2005), 685-689.

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TIMIS2012 141 s t Annual Meeting & Exhibition

Randall M. German Honorary Symposium on

Sintering and Powder-Based Materials

Edited by: K. Morsi

Fernand D.S. Marquis John L. Meyer

Ahmed EI-Desouky Eugene Olevsky

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materiah Society), 2012

A REVIEW ON ALLOYING IN TUNGSTEN HEAVY ALLOYS

Animesh Bose1, Rajendra Sadangi2, and Randall M. German3

'Materials Processing, Inc., 2421 Thomas Road, Haltom City, TX 76109 2Dept. of Materials Science and Engineering, Rutgers University, Pisctatway, NJ 08854 3College of Engineering, San Diego State University, 5500 Campanile Drive, San Diego,

CA 92182

Keywords: Tungsten Heavy Alloy, Liquid Phase Sintering, Contiguity, Mechanical Properties, Alloying, Powder Injection Molding

Abstract

This paper will review some of the developments in the area of alloying of tungsten heavy alloys. The review will concentrate on the alloying additions that have been made to the classic liquid phase sintered tungsten heavy alloys primarily based on W-Ni-Fe compositions. The effects of these alloying elements on the microstructure and properties of tungsten heavy alloys will be discussed in this paper. The review will also touch on some of the advantages that these additive modified tungsten heavy alloys yield over conventional tungsten heavy alloys (without the alloying additions), especially the ability to form near net-shaped, high strength, high hardness, heavy alloys without the additional thermc-mechanical treatments needed to attain higher strength and hardness.

Introduction

Tungsten heavy alloys (WHA) are two-phase composite which are one of the earliest materials fabricated by the powder metallurgy route [1-3]. The alloys exhibit a unique combination of high density, strength, ductility, and corrosion resistance. The alloys were developed primarily for high density uses, fabricated in a cost-effective manner. The consolidation of pure tungsten (density 19.3 g/cc) required excessively high temper-atures, which made the fabrication process expensive. The solution to the above problem evolved through the use of a mixture of tungsten powders with a small amount of lower-melting metal or combination of metals preferably with some solubility for tungsten. During sintering the low melting additive(s) forms liquid and the material would rapidly density to form a fully dense two-phase structure in a relatively short time. Thus, the additives enabled consolidation of tungsten at significant lower temperatures to fabricate the tungsten heavy alloys.

The lower melting elements were selected such that the liquid phase dissolves some tungsten into solution, inducing wetting of the tungsten grains, resulting in tungsten grain rearrangement and concomitant densification, resulting in the attainment of full density. Early on it was found that pure copper resulted in poor densification, even though the copper forms a liquid at the sintering temperatures. Copper had practi-cally no solubility for tungsten. It was also determined that the use of copper with

455

nickel could produce almost fully dense alloys as the liquid formed was able to wet the tungsten grains and take some tungsten into solution. Thus, early alloying additives consisted of nickel and copper which was generally sintered in hydrogen at temperatures around 1450°C. Fully dense tungsten alloys having a composition of 90W-6Ni-4Cu was fabricated in the late 1930s [3]. Another alloy having a composition of 90W-5Ni-5Cu alloy was also consolidated to a near full density which was around 16.5 g/cc (~ 50% more dense than lead). However, more recent tungsten heavy alloys consist of different alloying elements most of which have some solubility for tungsten. The combinations of transition elements such as nickel, iron, and cobalt have proven to be excellent alloying elements for fabricating modern day tungsten heavy alloys, with W-Ni-Fe-based alloys being the industry standard.

Conventionally, the needed strength and hardness are obtained by thermo-mechanical treatments after sintering (combinations of swaging and aging treatment). Unfortunately, this puts a limitation on the net-shape fabrication which is a desirable feature of powder metallurgy (PM). Extensive worldwide PM research has been directed to develop novel higher strength and higher hardness WHAs without the conventional thermo-mechanical treatments. Success in this approach is expected to lead to significant microstructural refinement with desirable flow characteristics, making these materials suitable for appli-cations in Kinetic Energy (KE) penetrators. Another benefit would be the ability to form complex net- or near-net shaped components by of powder injection molding methods, as the need to swage or work the material after sintering would no longer be necessary to produce the higher strength and hardness).

Before embarking a review on the alloying additives, it would be prudent to briefly outline uses and general properties of WHAs along with some of the drivers that have prompted research into novel WHA composites. However, it would be beyond the scope of this review to discuss alternate alloying and new matrix alloy compositions that have been attempted since the 1990s. Instead, this review will concentrate only on alloying of conventional W-Ni-X or W-Ni-Fe-X type alloys that have been investigated (where X is an alloying additive other than W, Ni, and Fe).

Genera l P r o p e r t i e s a n d Uses

Tungsten heavy alloys exhibit a unique combination of high density (typically >16 g/cc), high strength, moderate hardness, good ductility, moderate toughness, excellent radia-tion shielding capability, moderate electrical and thermal conductivity, vibration damp-ening ability and good corrosion resistance. The mechanical properties of WHAs can be tailored through compositional changes (influenced both by tungsten content and the Ni:Fe ratio), processing conditions (both during sintering and post-sinter treatments), and the final microstructure. Several research groups have investigated the effect of Ni:Fe ratio on the properties of WHAs and the compositions based on either 7Ni:3Fe or 8Ni-2Fe ratio have displayed optimum property combination [4-6].

456

In general, with increasing tungsten content (in the range of 88-94 wt % tungsten), the ultimate tensile strength, yield strength, and hardness of the material tends to in-crease while the elongation of the alloy decreases. In some special grades the tungsten content can be as high as 98 wt%. However, increasing the tungsten content above 94 weight percent can often result in a decrease in ductility and also ultimate tensile strength as the failure of the WHA is then dictated more by the microstructural char-acteristics of the two phase composite (a drastic increase in contiguity). Thus, it can be said that the ultimate tensile strength (UTS) of WHAs exhibits a maximum and then drops off with increasing volume fraction of tungsten while the ductility tends to show a continuous drop. For example, maintaining the same Ni:Fe ratio of 8:2, but changing the tungsten content from 90 to 95 to 98 (wt.%), the UTS changes from 925 to 955 to 708 MPa, respectively while the elongation continually decreases from 36% to 17% to 1.5%, respectively.

The processing parameters play a key role in determining the properties of WHAs. The use of hydrogen is necessary to reduce oxides and dissolved oxygen in the powders. However, use of hydrogen can lead to hydrogen embrittlement, which can be removed through the use of via vacuum annealing. This results in a significant improvement in the properties of the WHAs, especially the ductility as this prevents the segregation of trace impurities to the W-matrix boundaries [7]. Optimal properties can be obtained through the initial use of dry hydrogen atmosphere, followed by a wet hydrogen atmo-sphere (during liquid phase sintering) and a final nitrogen atmosphere to reduce the effect of hydrogen embrittlement [8]. The sintering temperature also influences the final properties of the WHAs. In general, in the liquid phase sintering range, an increase in the sintering temperature results in a decrease in the strength but an increase in the ductility of the material. The effect of increasing temperature manifests itself through modifications in the microstructure of the WHA (increase in W grain size as well as reduction in contiguity). A special post-sinter thermal cycling process has also been re-ported to improve the mechanical properties by lowering the contiguity in the tungsten phase and promoting strong tungsten-matrix interface [9].

It should be pointed out that variations in the processing conditions tend to affect the toughness of the WHAs. Ductility , strength and hardness are not that sensitive to process variations as toughness. Significant increases in strength and hardness in WHAs can be achieved through post-sinter thermo-mechanical processing of swaging and aging. The as-sintered WHAs subjected to a reduction in area (through swaging or rolling) followed by a slightly elevated temperature expasure (aging) for an extended period of time results in a significant increase in strength of the material. This process is often used to achieve high strength, high hardness WHAs [10-11].

The unique property combination that is exhibited by these two-phase tungsten al-loys has resulted in the use of this class of material in numerous applications [12-16]. The high density of the WHAs makes it an attractive candidate for KE penctrator and coun-terbalance related applications while the ability of tungsten to absorb radiation lends itself to its use in radiation shielding applications. Other possible applications include

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counterweights for aircrafts and helicopters, counterweights for self-winding wristwatches and various instruments, gyro-rotors, and radiation shields for shielding against gamma radiation, vibration dampeners (used as tool holders to reduce chatter during machin-ing), sporting goods (such as golf club heads), heavy duty electrical contact materials that can resist arcing, and weights used in numerous cell phone applications. Another important application Molybdenum containing grades is in case hot tooling dies for form-ing Titanium alloys. In case of KE penetrators, the candidate material that has been extensively used in the past, depleted uranium (DU) alloy, is being phased out due to environmental concerns (the material has been banned in Europe), while the known tox-icity of lead has led to its replacement by WHAs. Most of the WHA properties discussed above are primarily based on quasi-static properties. Due to the extensive application of this material in KE penetrators, an extensive study in the behavior of this alloy under dynamic conditions (high strain rate) as well as during ballistic penetration has been carried out. Though a detailed discussion of the dynamic and ballistic properties of WHAs would be outside the scope of this paper, it would still be important to provide a brief background about this area.

High density is a basic requirement for KE penetrators. WHAs can be fabricated to densities comparable to DU and quasi-static properties that are equal to better. However, the DU alloys outperform WHAs due to their ability to self-sharpen during penetration events [17]. This self-sharpening ability of DU during penetration events is attributed to the phenomena of adiabatic shear localization or flow softening [18-21]. Conventional WHA forms a mushroom-head during penetration and this results in a marked decrease in the penetration performance [22]. Thus, major efforts are underway world-wide to fabricate WHAs that will undergo flow softening and thus, promote self-sharpening be-havior instead of mushrooming. In simplistic terms, if the rate of hardening of a material undergoing plastic deformation is lower than the rate of softening caused by the conver-sion of work to heat, then the deformation will eventually become unstable. At that point, it is believed that the deformation would begin to localize. Among the pathways that have been investigated include new alloy systems containing metals that are prone to adiabatic shear band formation, increasing the hardness and strength and creating more homogeneous and finer microstructures (smaller tungsten grains). It was felt that if addition of certain elements could aid in the formation of finer microstructures in WHAs and also increase its strength and hardness, the material would be prone to self-sharpening [23,24]. This paper will review some of the alloying additions that have been investigated and their effects on microstructure and properties of conventional WHAs. To do that, a brief discussion about the WHA microstructure and its effects on properties would be helpful.

M i c r o s t r u c t u r e and Property

A typical WHA microstructure consists of almost pure bcc tungsten grains embedded in a relatively soft and ductile fee alloy matrix [25-31]. The WHAs are consolidated using the classic liquid phase sintering process and there are several stages of microstructural evolution. The first stage is solid-state sintering when numerous necks are formed be-

458

tween the powder particles. Then, as soon as the liquid is formed, the liquid wets the solid tungsten grains resulting in rapid particle rearrangement. The capillary force due to the wetting of the liquid is the driving force for rearrangement of tungsten grains. The liquid effectively penetrates loosely packed particle clusters and redistributes them. The next step is known as the solution-reprecipitation step. For this to happen, the solid has to have some solubility in the liquid phase. Since the transitional elements have some solubility for tungsten, the liquid also starts to take into solution some tungsten (dissolution of the finer particles are favored) and then reprecipitates tungsten on larger tungsten grains (that have not been dissolved). During this step there is significant grain shape accommodation as well as pore elimination. There is a decrease in the net surface energy due to pore elimination by the liquid. The final step in the process is microstructural coarsening with very little densification. During this last stage there is a drive to minimize the surface energy per volume ratio. The solution-reprecipitation event still continues. Grain growth occurs primarily by grain coalescence. Generally, grain coarsening follows a power law with time and can be written as: G"(t)-G"(o) = kt , where G(o) is the initial grain size, G(t) is the grain size at the isothermal sintering time t, and K is the growth rate constant (dependent on material parameters). The exponent n is generally 3 for solution-reprecipitation controlled event. Yang et. al. [32] took into account the effect of contiguity on the growth kinetics of liquid phase sintered W-Ni-Fe alloys and observed good agreement with experimental results.

As discussed earlier, the microstructure has a great influence of the properties of WHAs. Contiguity of tungsten and mean free path in the binder phase play important role in determining tensile strength and ductility. With increasing volume fraction (~ 0.75 - 0.95) of tungsten in the WHAs, the UTS initially increases and reaches a maximum around 0.85 volume fraction of tungsten and then starts to decrease. However, the elon-gation of the material decreases continually with increasing volume fraction of tungsten. Above a certain volume fraction of tungsten, contiguity and connectivity of tungsten in-creases and the fracture mode changes to intergranular failure. Under optimal processing conditions, W-W interface is the weakest link in the microstructure, which increases with increasing contiguity. Thus, an increase in the contiguity above a certain level results in the lowering of strength in the WHAs. A post-sinter heat treatment process, comprising of repeated thermal cycling around 1150°C has been used to decrease the contiguity of WHA microstructure concomitant improvement in the impact properties [23,28].

Alloying additions, such as molybdenum [33-36], tantalum [37,38], and rhenium [39,40] to WHAs have been quite effective in increasing the strength of the WHAs. The incorporation of these alloying additions also results in the refinement of the mi-crostructure. Among the alloying additives that have been used, rhenium seems to have the greatest effect on increasing the strength. In fact, ~ 2 wt.% addition of Rhenium can result in a hardness increase that can be achieved by > 6 wt.% addition of Mo and strength increase that can be matched by > 8 wt.% Mo addition. Figure 1 shows the microstructures of a 90W-7Ni-3Fe alloy, a 86W-4Mo-7Ni-3Fe alloy, a 88W-2Re-7Ni-3Fe alloy and a 85W-5Ta-7Ni-3Fe alloy, where the grain refinement with the alloying ad-ditives is clearly evident [40]. The increases in strength and hardness with increasing

459

alloying additions seemed to have a linear relationship at least in the range of 2 to 8 wt.% addition of the alloying element (replacing the tungsten in the alloy). Thus, two simple equations can be generated for Mo and Re additions [40]:

Yield Strength (MPa) = 524 + 1635 XMo where XMo is the atomic fraction of Mo. Yield Strength (MPa) = 563 + 5712 X^, where Xm is the atomic fraction of Re.

Figure 1. Representative microstructures of tungsten heavy alloys with different compositions (a) 90W-7Ni-3Fe, (b) 86W-4Mo-7Ni-3Fe, (c) 85W-5Ta-7Ni-3Fe, and (d)

88W-2Re-8Ni-2Fe[40]

Intermetallic compounds with Ll2-structure have been investigated as matrix phase for WHAs [41]. It is hypothesized that adiabatic shear can be enhanced if the yield strength of matrix phase is increased with temperature. It is well known that conventional metals and alloys loose strength with increase in temperature. On contrary, the Ll2-structured intermetallic compounds show an increased in strength up to a certain critical tempera-ture and then the strength eventually decreases. Novel tungsten heavy alloys containing nickel aluminide (Ni-25A1 at.%) and Ni-12Al-40Fe (at.%) alloy matrix showed promise of shear localization. However, they displayed lower room temperature strength and ductility.

Manganese was also used in the alloy matrix because it has a lower melting point, low thermal conductivity and comparable specific heat to iron. It is hypothesized that if the heat generated by the localized flow cannot be dissipated rapidly, shear bands can develop due to unstable internal strain gradients, characterized by very high local strain rates and near melt temperatures. Therefore, factors such as low thermal conduc-tivity, specific heat, density, strain hardening rate, high shear yield stress and thermal

460

softening rate are likely to promote adiabatic shear. On this basis, the W-Ni-Mn alloy matrix could be prone to adiabatic shear. Manganese also has a lower melting point than iron and a suitable Mn-Ni composition can significantly lower the matrix liquidus temperature, compared to other conventional heavy alloys. A 40Ni-60Mn alloy forms a liquid at 1014°C, which is much lower than the Ni-Fe alloy. A 90W-4Ni-6Mn alloy was liquid phase sintered at 1300°C to a density of 16.3 g/cc (~ 96% of theoretical). The microstructure of the sample contained porosity and unreduced oxides. Nevertheless the microstructure of the sample resembled that of conventional W-Ni-Fe heavy alloys, with rounded grains of tungsten dispersed in an alloy matrix. The average grain size of this new alloy was 7/itn, which is considerably finer than the grain size of conventional heavy alloys (20 to 50/rni). The matrix composition contained a small trace of tungsten (2.2W-43.3Ni-54.5Mn).In comparison, the matrix alloy of a conventional WHA shows a composition of 24W-53Ni-23Fe. The alloy exhibited significant shear localization. For the ease of processing, this alloy can be modified to use a higher Ni:Mn ratio. Higher Ni:Mn ratio of 3:2 WHA also proved to be brittle in the as sintered condition [44]. How-ever, the strength and ductility was increased to 900 MPa and 10% elongation through a post sinter heat treatment. The high strain rate properties of the material still needs to be investigated.

Another interesting alloying additive that has been examined is boron carbide (both as B4C as well as elemental boron and carbon addition). A new alloy made from 90W-6Ni-l.5Fe-l.96B-0.54C (wt.%) as well as 90W-6Ni-1.5Fe-2.5B4C (wt.%) have shown promise as tool bits as they possess significantly higher hot hardness [45]. Boron doping of WHAs have resulted in strength improvements. However, doping is not going to be covered in this paper.

Several other alloying additives have also been investigated by various research groups. Most of these developments have been directed to promote shear localization at high strain rates, similar to DU alloys. However, the use of alloying additives to increase the strength and hardness of such materials without thermo-mechanical treatments can lead to other applications that have complex geometries. For these applications with com-plex shapes, a near net shape processing technique of powder injection molding (PIM) of WHAs with alloying additives to increase strength and hardness can be utilized. PIM is a relatively new metal shaping technique that uses a powder-binder mixture that can be injection molded into complex shapes and then debound and sintered to achieve com-plex, net or near-net shapes [46]. Such as process has been successfully used to process conventional WHA [47] as well as high strength high hardness heavy alloys [40,48]. Table I shows typical properties of PIM tungsten heavy alloys.

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Table I: Properties of powder injection molded tungsten alloys and composites [40]

Composition (wt%)

Yield Ultimate Tensile Hardness Elongation Strength (MPa) Strength (MPa) (HRA) ( %)

90W-8Ni-2Fe Press and sinter 1500°C

82W-8Mo-8Ni-2Fe Press and sinter 1500°C

82W-8Mo-8Ni-2Fe Injection molded 1475°C

82W-8Mo-8Ni-2Fe Injection molded 1500°C

82W-8Mo-8Ni-2Fe Injection molded 1530°C

93W-5Ni-2Fe Press and sinter 18% reduction

551

688

775

700

688

918

1048

1144

1115

1067

1200

63.5

65.9

67.6

64.4

64.4

70.9

36

24

8

20

27

13

Future Directions

Tungsten heavy alloys are two-phase composites generally fabricated by sintering of mixed elemental powders. The structure comprises of interconnected tungsten grains with interpenetrating matrix phase. There is a significant growth of tungsten grains, which results in lower strength and hardness. Thermo-mechanical processing such as swaging and aging is commonly used to increase strength and hardness of WHAs. This paper has described an alternate way of fabricating high strength high hardness WHAs through the addition of additives such as Ta, Mo, Re, Ni-Al, Ni-Al-Fe, and B-C, which show promise for realizing finer grain tungsten heavy alloys with improved properties.

Microstructural evidences suggest that superior properties can be obtained if finer grain sizes can be retained in the sintered product. Higher hardness and transverse rupture strength have been achieved in submicron-grained cemented carbides, which is also a classic liquid phase sintered system. Submicron-grained carbides are routinely produced by careful control of raw materials and conventional sintering practice. It is envisioned that improved mixing of constituent metal powders and careful choice of alloying element(s) will be key to realize submicron-grained tungsten heavy alloys. Better densification characteristic is expected from finer powders with homogeneous distribution of constituents. This can lead to a microstructure with finer binder mean free path and low contiguity in tungsten phase. Such a structure is expected to promote adiabatic shear banding and hence, improve the performance of KE penetrators. Also, the use of powder injection molding in conjunction with these new additive modified WHAs results in the exciting possibility of being able to fabricate high strength high hardness WHAs into complex shapes. It is expected that the developments in the area of novel alloying additions to WHAs to improve the KE penetrator materials will continue while new applications will be developed for complex net-shaped WHA parts having higher strength and hardness.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

LOW-THERMAL-LOAD CONSOLIDATION OF SM-FE-N FLAKE POWDER BY COMBINATION OF CYCLIC COMPRESSION AND

CURRENT SINTERING

Kenta Takagi, Hiroyuki Nakayama, Kimihiro Ozaki

National Institute of Advanced Industrial Science and Technology (AIST): 2266-98, Anagahora, Shimoshidami, Moriyama-ku, Nagoya, 463-8560, Japan

Keywords: Current sintering, SmiFeyNx magnet, Flake-form powder, Densification

Abstract

Sm-Fe-N alloys, which possess a high magnet performance, are difficult to sinter due to their thermal decomposability. SmiFeîN magnets were produced from flake powders by cyclic high-pressure compaction and subsequent current sintering with low temperatures. Under the high pressure compression, the flake particles were densely packed while being broken up and orderly stacked. Coarser powder resulted in denser compacts. When the coarse powder of < 355um was used, it was densified beyond 85% in relative density only by the cyclic compaction with the pressures above 1.2GPa. The compacts were rigidly consolidated by the current sintering in the temperature range of 350~400 CC without the decomposition. Transmission electron microscope (TEM) observation verified interparticle sinter-bonding in the compacts. The developed process finally produced the sintered magnets with the high density up to 93%.

Introduction

Sm-Fe-N-based alloys are well-known as one of rare-earth-transition-metal hard magnet alloys with high magnetic properties; in particular, those with TbCu7- and Th2Zni7-type crystal structures exhibit an excellent magnet performance [1, 2], In addition to their high magnetic properties, these materials have a number of advantages: high Curie temperature, less coercivity -degradation at a high temperature and high corrosion resistance. Therefore, Sm-Fe-N magnets are preferred for applications in the motors exposed to a high temperature, such as hybrid and electric vehicles. Furthermore, due to price escalation and supply instability of heavy rare metal resources, Sm-Fe-N magnets that do not require these elements have been target as potentially alternative magnets to Dy-doped Nd-Fe-B magnets. Nevertheless, Sm-Fe-N alloys remain minor magnet material because they are applied only to bond magnets due to the difficulty in sintering. The bond magnets, which are composed of magnet powder and plastic binder, cannot bring out the high thermal resistance of Sm-Fe-N alloys.

The primary reason that the Sm-Fe-N alloys are difficult to sinter is their easy decomposition. These alloys gradually decompose into Sm-N and a-Fe above about 500 °C. Therefore, once Sm-Fe-N alloy powder was sintered above this temperature, its magnetic properties, especially coercivity, remarkably degrade [3]. Considering that the melting point of Sm-Fe alloys is around 1300 °C, temperatures below the decomposition point are too low to density powders of these alloys by solid-state sintering. Hence, many efforts have been made to sinter Sm-Fe-N powders using spark plasma sintering [4]. However, they have not led to the simultaneous achievement of densification and avoidance of coercivity degradation.

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We have so far developed a high-pressure current sintering technique for low-thermal-load consolidation of amorphous powders. This technique was believed to be exploitable as an underlying technique to densely consolidate Sm-Fe-N powders below the decomposition temperature. In Ulis method, densification is achieved by the high pressure compression and the applied current mainly serves interparticle sintering bonding. Therefore, it is important to design the consolidation process with a focus on the powder compaction. In general, use of fine powder is generally effective for a low temperature sintering. However, focusing on the powder compaction of the brittle material, coarse powder is rather effective to obtain dense compacts; coarse powder is closely packed under compression with breakup and rearrangement [5]. Therefore, we chose coarse powders with flake form as a raw powder. Flake powders were more easily broken up by the compression than granular powders, and furthermore they were expected to be densely packed with layer stack of particles. Besides the powder form, cyclic compression is known to have an effect on densification in powder compaction process. Thus, this study aimed to develop the low-thermal-load consolidation process for dense Sm-Fe-N bulk magnets by combination of the three techniques, the high-pressure current sintering, the cyclic compression and the flake powders. In particular, we investigated the effects of powder size distribution, cyclic compression condition, and sintering temperature oh densification. Here, in this study, Sm-Fe-N alloy with TbCu7-type crystal structures, i.e. SmiFe7Nx alloy, was chosen as a composition of the raw powder.

Experimental Procedure

Polydisperse SmiFe7N„ flake powder with a thickness of about 20 um was firstly prepared as a raw powder. This powder consisted of nanosize-grained polycrystalline because it was produced via a melt spinning. From this original powder, we prepared four types of powders (A~D) with different size distributions by a sieving to investigate the effect of size distribution, as shown in Fig. 1. These powders were put in a cylindrical mold set with an inner diameter of 6 mm and vibrated to induce close packing. The mold and punches were made of WC-Co cemented carbide to withstand the high pressure in the order of gigapascals. The mold containing the powder was set in an electric current sintering system equipped with a servo-type high-pressure unit (Plasman CSP-I, SS Alloy Co. Ltd., Japan). Prior to sintering, the powder underwent either single or cyclic uniaxial-compression under pressure of -1.8 GPa at R.T. In the case of cyclic compression, the pulsated pressure, which varied from 0 to the maximum pressure, was applied for up to 100 cycles. After the compression process, the compact was sintered in vacuum under the high pressure of 1.2 GPa. The sintering temperature was varied from 200 to 500 °C, which was lower than the decomposition temperature of SmiFeyN* alloy. To reduce the thermal load on the compact, the heating rate and holding time were set to 40 °C/min and 60 sec, respectively. Voltage and current applied to the compacts were 4 V and 800 A at a maximum.

The densities of the compressed and sintered compacts were determined by the calculation using the compact height in the mold or the Archimedes method. Phase identification and microstructural characterization of the sintered compacts were conducted using XRD, SEM and TEM. In order to verify the effect of sintering, the strengths of the compacted and sintered samples were simply evaluated by the Brazilian test [6] using a universal testing machine. Magnetic measurements were conducted using a pulse B-H tracer (PBH-1000, Nihon Denjisokki Co., Japan).

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Figure 1. Particle size distributions of the powders used in this study.

Figure 2. Densification behavior of Sm-Fe-N flake powder (powder C) during high-pressure current sintering.

Results and Discussion

Figure 2 shows the densification behavior of powder C during high-pressure current sintering at 400 °C. Assuming that this compact was fully densified, the height of the compact was calculated to be 3.7 mm. As predicted, heating at such a low temperature gave rise to the slight densification by only several percent, and much of densification occurred at the stage of the high-pressure compression. Table I shows the density, mechanical strength and magnetic properties for the compacts sintered at the various sintering temperatures. The density gradually increased with the sintering temperature in the range below 400 °C, and rose substantially over 450 °C. The compact sintered at 500 °C exhibited very high density over 96%. However, the coersivity began to decrease from 450 °C, and eventually it was dropped by one third of the raw powder at 500 °C. This decrease in coercivity would be caused by the thermal decomposition. In fact, small peaks derived from the a-Fe phase were found in the XRD profiles for the compacts sintered at 450 and 500 °C. These results suggest that the sintering temperature should be confined below 400 °C.

On the other hand, the diametral tensile strength (DTS) of the sintered compacts showed different behavior from the relative density, as indicated in Table I. Whereas the significant increase in density appeared at the sintering temperature above 450 °C, DTS began to increase from 350 °C in spite of the small increase in density. The increase in the DTS implied the occurrence of sintering bonding between the particles. The SEM cross-section of the compact sintered at 400 °C is shown in Fig. 3. The particles with flake form were stacked in layer as expected, and the adjacent particles were well bonded to such an extent that a large part of interfaces could not be identified. Indeed, the TEM observation demonstrated that the interfaces were bonded on atomic level. Consequently, the current sintering should be conducted in the temperature range from 350 to 400 °C to sinter the SmiFe7Nx powder without the thermal decomposition. This temperature range allows the interparticle bonding, but is not enough to

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promote the densification. Hence, densification must be previously achieved at the stage of high-pressure compression.

Table I. Various properties of the compacts sintered from the powder C at different temperatures. The applied pressure was fixed at 1.2GPa

Figure 3. SEM cross-section of the compact sintered from the powder C at 400 °C under 1.2 GPa.

In general, major influential factors for densification in powder compacting are applied pressure, size distribution, and process temperature. An enormous number of studies examining such subjects has been conducted down the ages, but the studies on flake powders are scarce. Fig. 4 indicates the changes in density with cycle number for the compacts pressed with the different pressure. The final density was increased with the pressure, and this increase seemed proportional to the pressure. In the beginning, we predicted that the compression energy given by the pressure over 1.2 GPa might be consumed for elastic deformations of the powder and mold based on the previous result of densification behavior obtained from a single-stroke compacting test. In other words, the flake powder is orderly and densely packed up to 1.2 GPa as seen in Fig. 4, and thus, densification had been thought not to evolve above 1.5 GPa. However, in the case of cyclic compression, the density was increased with the increase in pressure. This is probably because the cyclic compression continually induced the breakup and rearrangement of the powders. Consequently, it was demonstrated that the larger pressure produced the denser compact, and the improvement of 6% was obtained at a maximum compared with the single-stroke pressing.

Figure 5 indicates the relationship between the particle size distribution and relative density for the compacts that were pressed at 1.2 GPa for 10 cycles. As initially expected, the coarser powder could be more densely compacted. The result also revealed that the powder with the narrow size distribution was more highly densified than the powders with broad distribution.

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According to the general packing theory, powders with broad size distribution allow dense compactions. The reason why the flake powders showed the opposite phenomenon is probably because the fine particles embedded between the stacking large particles. Anyhow, in the case of flake powder, coarse powder with narrow size distribution was of advantage to densification.

Figure 4. Changes in relative density as a function of cycle number for the compactions applied with the different compressive pressure. The powder C was used.

Figure 5. Relationship between the particle size distribution and the relative density after 10-cycle compression with the pressure of 1.2GPa.

Finally, we tried to fabricate dense SmiFe7Nx bulk magnets using the optimized conditions of the cyclic compression and current sintering. Firstly, the powder C was chosen in this study, even though the powder D has been demonstrated to be preferred. This was because the preparation of the latter powder involved a large yield loss in sieving. The pressure applied in the cyclic compaction was determined to be below 1.2 and 1.5 GPa. According to the results described above, the higher pressure, 1.8 GPa might produce the denser sintered magnets, but this high pressure had the risk to break the mold and punches during processing at the present moment. The sintering temperature was fixed to 400 CC for the avoidance of thermal decomposition from the results of the examination of current sintering.

Table II shows the various properties of the resulting sintered compacts. The sample sintered via the cyclic compaction at 1.2 GPa exhibited a density of 90.6%, which was approximately 4%

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denser than the sintered samples derived from the single-stroke compression as indicated in Table I. Furthermore, the cyclic compaction at the higher pressure of 1.5 GPa resulted in the densest bulk with a density above 92% after sintering. Moreover, this sintered magnet exhibited a (BH)max of more than 16.2 MGOe. This value is much higher than the (BH)max of commercial isotropic Sm-Fe-N bond magnets (13 MGOe) and is close to the (BH)max value of anisotropic Sm-Fe-N bond magnets (18 MGOe) [7]. The obtained bulk magnets are expected to be applicable as an isotopic magnet subjected to hot environments because they have no organic bonding medium.

Table II. Density and magnetic properties for the sintered magnets which were produced from the powder C by the cyclic compaction and current sintering.

Compression condition Sintering Cycle Pressure (GPa) condition

100 1.2 1.2 OPa

400 "C 1 0 0 1 J for<S0 sec

Raw powder

Density Br Ha BHmm (%) (kQ) (kOe) (MGOe)

90.8 9.00 9.70 15.7

92.4 9.10 9.68 16.2

— 9.67 9.76 1S.5

Summary

Development of a low-thermal-load consolidation technique comprising the combination of high-pressure cyclic compaction and current sintering was conducted to produce dense bulk magnets of Sm-Fe-N, which has low thermal stability. The coarse SmiFe7Nx powders with flake form were used. Firstly, the single-stroke pressed compacts were sintered at the various temperatures. This result suggested that the sintering temperature must be confined in the range of 350 ~ 400 °C to avoid the thermal decomposition. Meanwhile, in the cyclic compacting, the higher compressive pressure was the more advantageous to densification. In addition, it was revealed that the raw powder should be coarser and have a narrow size distribution to obtain dense compacts. The sintered magnets which were produced based on the obtained knowledge exhibited the high relative density of more than 90%. Therefore, these magnets had the high {BH)max over 16 MGOe, which was much higher than that for commercial Sm-Fe-N isotropic bond magnets.

References

1. S. Sakurada, A. Tsutai, T. Hirai, Y. Yanagida, M. Sahashi, S. Abe, T. Kaneko, "Structural and magnetic properties of rapidly quenched (R, Zr)(Fe, Co)ioNx," J. Appl. Phys., 79 (1996), 4611-4613. 2. T. Iriyama, K. Kobayashi, N. Imaoka, T. Fukuda, H. Kato, Y. Nakagawa, "Effect of nitrogen content om magnetic properties of Sn^FenNx," IEEE Trans. Magn. 28 (1992), 2326-2331.

| 3. F.A.O Cabrai, S. Gama, E. de Morais, N.-L. Sanjurjo, C.A. Rubeiro, C.C. Colucci, "Study of thermal decomposition mechanism of the FenSn^Ns phase," IEEE Trans. Magn., 32 (1996), 4365-4367. 4. D.T. Zhang, M. Yue, X. Zhang, "Study on bulk Sm2Fei7Nx sintered magnets prepared by spark plasma sintering," Powder Metal., 50 (2007), 215-218. 5. R.J. Roberts, R.C. Rowe, "Brittle-ductile transitions die compaction of sodium chloride," Chem. Eng. Sei., 44 (1989), 1647-1651. 6. C. Pittet, J. Lemar, "Mechanical characterization of brushite cements,' 'J.Bio med. Mater.Res.,

| 53 (2000),.769-780.

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7. K. Ohmori, T. Ishikawa, "Progress of Sm-Fe-N anisotropic magnets," Proc. 19 Int. Workshop on Rare Earth Permanent Magnets & Their Applications, Beijing, China, Aug. 30, 2006, p. 221-230.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

FABICATION OF TIN / FE-AL CERMET FROM MIXTURE OF TIN, FE AND AL POWDERS

Hiroyuki Nakayama, Kimihiro Ozaki, and Keizo Kobayshi

National Institute of Advanced Industrial Science and Technology; Moriyama-ku Anagahora Shimoshidami 463-8560; Nagoya, 2266-98, Japan

Keywords: microstructure, bending strength, hard material, intermetallic compound

Abstract

TiN shows high hardness, good thermal stability and conductivity. Hence, fabrication of TiN cermet considering an application for cutting tools is desired. Fe-Al intermetallic compounds are one of a candidate for binder of the cermet, because it shows good mechanical property and it is composed of common (low cost) elements. Therefore, in this study, fabrication of TiN / Fe-Al cermets were examined.

TiN, Fe and Al powders were mechanically milled using a planetary ball mill under Ar atmosphere. The nominal composition of powder mixture was TiN-10 mass% FeAl. The milled powder was consolidated by pulsed current sintering. The sintered compact was composed of TiN and FeAl intermetallic compound. The intermetallic compound would be formed by the reaction of molten Al and solid Fe during the sintering process. The milling conditions affected the three point bending strength, which showed 0.96 GPa using the appropriate milling conditions.

Introduction

Titanium nitride (TiN) exhibits some superior properties, such as high hardness, good thermal / chemical stability and thermal conductivity [1], but a brittle and low sinterability. Therefore, a consolidation with metal binder (cermet) is desired. However, the wettability of the TiN and pure metal is quite low, the wetting angles with Fe, Co and Ni are over 90 degree. In addition, long period liquid-phase sintering led to a partial decomposition of the TiN [2], In the previous research, the process of mechanical milling of TiN and Fe and subsequent pulsed current sintering (PCS) of the milled powder can fabricate a dense TiN/Fe cermet showing a three point bending strength of 1 GPa [3]. However, in the usage of Fe as a binder phase, gathering rust is inevitable, so that the replacement of a binder phase which exhibits high oxidation resistance is favorable. One of the candidates for the binder phase is Fe-Al intermetallic compound, because the compound composed a common elements and good oxidation resistance. In the actual case, cemented carbide with FeAl binder (WC/FeAl) has shown the higher oxidation resistance compared to the conventional cemented carbide (WC/Co) [4]. Therefore, in this study, the fabrication of TiN/Fe-Al cermets was examined using mechanical milling and PCS process.

Experimental Procedure

The cermet with TiN-90(Fe-Al) mass % composition was fabricated. The mechanical millings were carried out by 2 steps. In the first milling, the TiN and Fe powder were

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mechanically milled for 14.4 or 28.8 ks using a planetary ball mill (Fritsch P-5) under Ar atmosphere. After the first milling, Al powder was added into the TiN and Fe mixture, and then second milling was carried out. The rotation speed was set to 2.17 times to the revolution speed. The rotation direction was opposite to the revolution. For the milling, chromium steel vials with 500 ml volume and cemented carbide balls with 5 mm in diameter were used. The ball to powder weight ratio was 0.07 (ball : 600g, powders : 40 g). The milled powders were consolidated by PCS method. The powder filled into a carbon mold with outer and inner diameter of 50 mm and 30 mm, respectively. The filled powder was pressed using a carbon punches at 60 MPa, and then sintered at 1473 K or 1573 K for 300 s. In the table 1, the detailed milling and sintering conditions used in this study are listed, and here after, the condition number is used as a sample name. Microstructure of the sample was observed by SEM. For the SEM observation, the sample surface was polished using a diamond paste with 1 p.m. Three point bending test with a span of 10 mm was performed. For the bending test, the samples were cut to 3.5 x 4.0 x 20 mm dimensions. The bending test repeated 5 times and the average value was used for the strength.

Table I. Condition of sample preparation.

Results and Discussion

Fig. 1 shows the SEM images of source powders of (a) TiN, (b) Fe, (c) Al and (d) milled powders of sample 1. The TiN powder has angular shape and the particle size widely distributed ( 1 p.m - 20 u,m). The Fe powder shows spherical morphology and the size of particles are homogeneous. The larger particles size compared to the other particles and awkward shape is

Figure 1. SEM images of source powders (a) TiN, (b) Fe, (c) Al and (d)milled powder of sample 1.

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observed in Al powder. After the milling, the particles significantly refined, and the aggregation of powders is seen.

The XRD result of the powder and sintered sample using the condition 1 is shown in Fig. 2. The sintered compact is composed of TiN and Fe-Al intermetallic compound. The Fe-Al compound was not observed in the milled powder, the peaks for source powders (TiN, Fe and Al) were seen, thus the Fe-Al formed during the sintering process by the reaction between Fe and molten Al. Other samples also showed the same result. Namely, the peaks for TiN and Fe-Al were detected after the sintering.

Figure 2. XRD patterns of as-milled powder and sintered sample.

The microstructure of sample using the condition 1 shows in Fig. 3 (a). There are three regions, white, gray and black, are seen. According to the EDX analysis, the large white regions showed the condensation of Fe and less of Ti are observed. In contrast, the black region is enriching Al and O content. Therefore, the white regions seem to be a Fe-rich Fe-Al or Fe phase. The black region would be an aluminum oxide. However, the Fe and oxide phase could not be detected in the XRD analysis, because of the small diffraction intensities of these phases. The other gray regions were Ti-rich phase, thus these were expected to be TiN. To suppression of the aluminum oxide formation, carbon powder as a reduction agent was added in the milling process. Fig. 3 (b) shows the SEM image of the sample 4. It is clear that the total amount of the black region reduces as compared with the sample without C addition. Hence, the C addition inhibits

Figure 3. SEM images of sintered (a) sample 1 and (b) sample 4. Sample numbers refer to Table I.

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the aluminum oxide formation. Fig. 4 shows the summary of the three point bending test. The sample 1 shows the 0.71 GPa.

The sample 2 shows the smallest strength of 0.31 GPa in this study. In this sample, the milling time of first milling was longer than that of the other samples, so that the over milling led to the degradation of the sample strength. The strength of the sample 3 exhibits the highest average value (0.96 GPa) in this study. In this case, the milling condition in the second milling was milder than that of the sample 1 and 2. The revolution speed of the sample 1 and 2 were 200 rpm. In contrast, the revolution of the sample 3 was 170 rpm. The total milling energy depends upon the milling time and speed, thus the above results implies the appropriate milling energy exists for realizing a high strength. The carbon added sample (sample 4) shows the 0.75 GPa in the strength, even though the amount of the aluminum oxide was smaller than other samples. In this sample shows the wide strength distribution. The minimum strength in this sample is 0.30 GPa, and the maximum is 1.27 GPa. Due to the microstructural observation, the aluminum oxide distributed is homogeneous in the sample without carbon addition, but the oxide exists inhomogeneously in the C added sample. Therefore, there is a possibility that the inhomogeneous microstructure caused the degradation of strength rather than that of the sample without C addition.

Figure 4. Summary of three point bending test in the sample prepared by several conditions.

Conclusion

TiN-10 mass% FeAl cermet alloys were fabricated by the mechanical milling of TiN, Fe and Al powders and subsequent PCS method. In the microstructure of the sintered sample, aluminum oxide formation was observed. Theses oxide formation was suppressed by C addition as a reduction agent in the milled powder. The milling conditions strongly affect to the three point bending strength in the sample. That is, the bending strength exhibited the trend that the longer time or higher energy milling degraded the strength. The controlling the milling condition achieved the relatively high bending strength of 0.96 GPa.

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References

1. H. O. Pierson, Handbook of Refractory Carbides and Nitrides, Westwood, NJ: Noyes Publications, (1996).

2. J. -K Park and S. -T Park, "Densification of TiN-Ni cermets by improving wettability of liquid nickel on TiN grain surface with addition of M02C", International Journal of Refractory Metals andHardmaterials, 17 (1999), 295-298.

3. H. Nakayama, K. Kobayashi and K. Kikuchi, "Characteristics of TiN-Fe Cermet Fabricated by Mechanical Milling and Pulse Current Sintering" Supplemental proceedings: Materials Processing and Energy Materials TMS, 1 (2011), 505-509

4. R. Subramanian and J. H. Schneibel, "FeAl-TiC and FeAl-WC composites - melt infiltration processing, microstructure and mechanical properties", Materials Science and Engineering ,4 244(1998), 103-112.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

TRANSPARENT POLYCRYSTALLINE ALUMINA OBTAINED BY SPS: SINGLE AND DOUBLE DOPING EFFECT

Burcu Apak1, Halide Esra Kanbur1, Esra Ozkan Zayim2, Gultekin Goller1, Onuralp Yucel1, Filiz Cinar Sahin1

'Istanbul Technical University, Metallurgy and Materials Engineering Department 34469 Maslak, Istanbul, Turkey.

Istanbul Technical University, Physics Engineering Department 34469 Maslak, Istanbul, Turkey.

Keywords: Spark Plasma Sintering, Alumina, Doping

Abstract

Commercial nanocrystalline alpha alumina powders were used for fabrication of dense and transparent alumina by spark plasma sintering (SPS). High purity dopants such as MgO, Y2O3, TiÛ2 and CaO with 150 ppm amount were added to alumina and powder mixtures were densified by SPS between 1175 and 1300 °C using 80 MPa pressure for 5 to 20 min durations. The influences of MgO, Y2O3, ÜO2 and CaO single and double doping on density, hardness, fracture toughness and microstructures of alumina samples are investigated. The highest hardness value was measured as 24.2 GPa in the sample doped with 150 ppm MgO and 150 ppm CaO sintered at 1175 °C under 80 MPa pressure for 20 min. The fracture toughness values were ranged between 3.3 and 4.9 MPa.m"2. Alumina ceramic with 150 ppm MgO which was spark plasma sintered at 1175 °C for 20 min showed the highest Real Inline Transmittance (RIT) value as 55.5 %.

Introduction

Sintered aluminium oxide (AI2O3) is one of the most significant high technology ceramics with its high strength, hardness, thermal stability, excellent dielectric properties, corrosion and wear resistance at elevated temperatures. Alumina finds lots of application areas such as high voltage insulators, high temperature electrical insulators, grinding media, ballistic armour, wear pads and orthopedic implants with its mentioned properties. Also in recent years, alumina is considerably attracting attention because of its transparent characteristic [1-3].

Transparent materials can be produced as crystalline or polycrystalline structures. When it is compared with mono crystalline structures, polycrystalline transparent ceramic materials have lower manufacturing costs and it is easier to obtain products with large diameter. With their excellent mechanical properties and low production costs polycrystalline transparent ceramics are gradually sustaining more importance in industrial applications. However, sintered polycrystalline ceramic materials show microstructural features like grains, grain boundaries, second phases and pores that significantly affect optical properties.

In polycrystalline materials translucency/transparency is determined with the amount of scattering light. Incoming light is scattered with the inclusions, pores, grain boundaries, phase boundaries and other surface or internal faults that are present in microstructure, so the microstructure of polycrystalline ceramics plays an important role in light transmission [2].

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Because of ease in accessibility, alumina is preferred ceramic for researchers. AI2O3 can be defined as a basic ceramic for sintering techniques to understand the sintering mechanism for novel routes such as microwave sintering [4, 5], solar sintering [6], spark plasma sintering [7-13] etc. Spark plasma sintering (SPS) is a consolidation technique that can combine high heating and cooling rates with an uniaxial applied pressure resulting in short processing time. The first unique property of the SPS process is the possibility of using very fast heating rates and very short holding times (minutes) to obtain fully dense samples. Also, when compared with the samples sintered with using conventional densification techniques, spark plasma sintered samples are usually reported to have been produced at lower temperatures. In SPS, heating is achieved through a dc pulsed current passed through the die as well as the sample. Normally, as an external heating element, graphite die and punches are used. In addition, part of the pulsed current passing through the sample results in sparking at interparticle contacts and leads to internal heating of the sample [7].

To modify the microstructure and mechanical properties of alumina ceramics, MgO,' Y2O3, TiÛ2 and CaO are widely used as dopant materials [7,14-22]. During sintering of alumina, MgO doping suppresses the grain growth and influences the surface diffusion [14-15], The grain boundaries do not break away from the pores, which prevents the inclusion of pores trapped inside new large grains, with slow/long diffusion path densification. The mechanism by which MgO slowdown grain boundary movements can be linked to the fact that the majority of MgO doped into AI2O3 resides at the grain boundaries, due to the relatively large difference in ionic radius, 0.72 À for Mg2+ and 0.53 À for Al3+. Yttria is an important dopant for increasing creep resistance of alumina and hindering grain growth during sintering [16-19]. Yttrium decreases the coarsening rate relative to densification rate and hence shifted the grain size-density trajectory to higher density for a given grain size. TiÛ2 has a beneficial role on densification but it promotes grain growth [20].

In the present study, the effects of single and co- doping on microstructural and mechanical properties of spark plasma sintered alumina ceramics are investigated.

Experimental Procedure

Commercial a-alumina powders (TM-DAR, Taimei Chemicals Co. Ltd., Japan) with an average particle size of 100 nm and a purity of 99.99%, were used in the present study. Doping elements were introduced into alumina before calcinations as Mg(N03)2-6H20, Y(N03)3-6H20, TiÛ2 and CaCl2.2H20. Batches were prepared by mixing Merck quality dopants and alumina in ethanol medium by ball milling for 24 hours with alumina balls. The slurry was then dried. The decomposition was carried out at 750 °C for lh.

Without any binder addition or pre-shaping application, the powder was directly loaded in a graphite die with 50 mm in inner diameter and sintering was carried out by using the SPS apparatus (SPS-7.40MK-VII, SPS Syntex Inc.). After applying a pressure of 80 MPa, low heating regime was used for sintering: 100 °C/min from room temperature to 600 °C, 25 °C/min from 600 °C to 1000 °C followed by 8 °C/min to sintering temperature (1175 or 1300 °C). The temperature of the die was measured by an optical pyrometer. The samples sintered at 1175 °C were held 20 min at the sintering temperature, whereas 5 min soaking time was applied to the sample sintered at 1300 °C. Samples heated with low heating rate (HR) were annealed at 1000 CC for 10 min subsequently. Whole process were carried out in vacuum atmosphere with recording shrinkage, displacement, heating current, and voltage for every 5

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sec. At the end of process, sintered disks with diameter of 50 mm and thickness of 15 mm are obtained. After sintering process, Archimedes method was used to determine the final densities of the compacts. Specimens, polished with a diamond paste having particle size of 1 um, were thermally etched at 1100 °C for 1 hour. The hardness and fracture toughness of the samples at room temperature were evaluated by the Vickers indentation technique at a load of 1 kg. The micrographs of all sample surfaces were observed by scanning electron microscopy (SEM; Model JSM 7000F, JEOL, Tokyo, Japan). The in-line transmittance was measured in the wavelength range 0.19-1.1 um using a spectrophotometer (Agilent, UV-Vis Spectrophotometer).

Results and Discussion

Density values of spark plasma sintered consolidated ceramics at 1175 °C for 20 min and 1300 °C for 5 min applying 80 MPa pressure both heated with 8 °C/min) under vacuum atmosphere are given in Table 1.

Table 1 . Relative density values of spark plasma sintered alumina ceramics Dopants

MgO

Y203

Ti02

CaO

MgO + Y203

MgO + Ti02

MgO + CaO

Sintering Procedure 1175 °C for 20 min 1300 °C for 5 min 1175 °C for 20 min 1300 °C for 5 min 1175 °C for 20 min 1300 °C for 5 min 1175 °C for 20 min 1300 °C for 5 min 1175°Cfor20min 1300 °C for 5 min 1175 °C for 20 min 1300 °C for 5 min 1175°Cfor20min 1300 °C for 5 min

Relative Density (%) 99.9 99.9 99.8 99.7 99.5 99.4 99.2 99.6 100 100 99.7 99.9 99.5 99.8

Relative density values were close to each other changing between 99.2 and 99.9 % in single doped alumina ceramics. As can be seen from Table 1, spark plasma sintering of alumina doped with 150 ppm MgO and 150 ppm Y2O3 resulted in higher relative density values compared to Ti02 and CaO additions with same amount. Increasing the soaking time to 20 min, obtaining such high relative density values is possible even in lower sintering temperatures such as 1175 °C. Co-doped alumina samples resulted in higher density values rather than single doped alumina ceramics. With the addition of MgO and Y2O3 together, theoretical density values were obtained both at 1175 and 1300 °C spark plasma sintering temperatures. When relative density results of co-doped alumina ceramics are considered, soaking time is a more effective parameter than spark plasma sintering temperature in case of removing the pores in the structure.

The hardness values of single doped alumina ceramics altered between 21.7 and 24.7 GPa. The highest hardness value was attained in the sample with 150 ppm CaO spark plasma sintered at 1175 °C for 20 min under applying pressure of 80 MPa. Increasing the spark plasma sintering temperature to 1300 °C and lowering the soaking time to 5 min usually

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resulted in decreasing hardness values. In co-doped samples, the hardness values were ranged from 21.4 to 24.2 GPa. In both single and co-doped samples, doping 150 ppm CaO to alumina ceramics increased the hardness results, but decreased the fracture toughness values. The highest fracture toughness as 4.5 MPa. m"2 is obtained in the sample containing 150 ppm Y2O3 spark plasma sintered at 1300 °C for 5 min with very low heating rates under applying pressure of 80 MPa.

Figure 1. Microstructures of alumina ceramics containing a. MgO b. Y2O3 c. TiC>2 d. CaO spark plasma sintered at 1175 °C for 20 min under a pressure of 80 MPa in a vacuum

Figure 2. Microstructures of alumina ceramics containing a. MgO + Y2O3 b. MgO + Ti02 c. MgO + CaO spark plasma sintered at 1175 °C for 20 min under a pressure of 80 MPa in a

vacuum atmosphere.

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Etched surface microstructures of alumina ceramics are illustrated in Fig. 1 and Fig. 2. Spark plasma sintering at 1175 °C for 20 min with 8 °C/min heating rate resulted in homogenous microstructures and small grains in both of the samples. Co-doping resulted in larger grain size compared to single doping. More homogenous structures were attained in single doped alumina ceramics. Residual pores in microstructure of alumina doped with 150 ppm Ti02 were removed in alumina ceramic co-doped with 150 ppm MgO and 150 ppm ÜO2.

The in-line transmission values of alumina ceramics with 0.7 mm thickness are represented in Fig. 3. 150 ppm MgO doped and spark plasma sintered at 1175 °C for 20 min under a pressure of 80 MPa showed highest inline transmission value in visible wavelength (400 - 700 nm). Whereas the highest in-line transmission value is reached in 470 nm wavelength as 66.6 %, the general in-line transmission ranged between 50-60 % and increased with increasing wavelength. The small grain size (~300 nm) and homogenous microstructure led to high transmission values.

Figure 3. In-line transmission of alumina ceramics versus wavelength

Increasing the SPS temperature to 1300 °C and decreasing the soaking time to 5 min caused a sharp decrease in in-line transmittance of alumina doped with 150 ppm MgO. In case of samples containing 150 ppm MgO and 150 ppm Ti02, spark plasma sintering at 1175 °C for 20 min resulted in higher transmission whereas, spark plasma sintering at 1300 °C for 5 min heating with low heating rates improved in-line transmission in 150 ppm Y2O3 doped alumina. CaO is the dopant affected alumina most negatively among alumina ceramics with four dopants (MgO, Y2O3, TiÛ2 and CaO) in both SPS temperatures. Very poor in-line transmissions were attained in the samples containing 150 ppm CaO. Co-doping affected the in-line transmissions of alumina ceramics negatively.

Conclusion

In this study, highly dense alumina ceramics with density values higher than 99.2 % were attained by spark plasma sintering at two different temperatures: 1175 °C and 1300 °C. 150 ppm MgO addition enhanced the densification mostly among four dopants (MgO, Y203> Ti02 and CaO). Co-doping MgO and Y203 resulted in theoretical density values. Vickers

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hardness of the samples ranged from 21.4 to 24.7 GPa, fracture toughness values of single and co-doped alumina ceramics were between 3.3 and 4.5 MPa.m . Because microstructures with small grains and without pores are needed for transparency, alumina ceramic with 150 ppm MgO which was spark plasma sintered at 1175 °C for 20 min showed the highest in-line transmission value. Increasing spark plasma sintering temperature to 1300 °C and decreasing the soaking time to 5 min resulted in increasing the in-line transmittance in case of 150 ppm Y2O3 doped alumina. Lower in-line transmission results were obtained in co-doped ceramics.

References

1. R. Klement et al., "Transparent armour materials", Journal of the European Ceramic Society, 28 (2008), 1091-1095. 2. S. Ma et al., "Sintering of translucent alumina," Journal of Materials Processing Technologies, 209 (2009), 4711^*715. 3. G.C. Wei, "Transparent Ceramics for Lighting," Journal of the European Ceramic Society 29 (2) (SI) (2009), 237-244. 4. J. Cheng et al., "Microwave sintering of transparent alumina," Materials Letters, 56 (2002), 587- 592. 5. Y. Fang, J. Cheng, and D. K. Agrawal, "Effect of powder reactivity on microwave sintering of alumina," Materials Letters, 58 (2004), 498- 501. 6. R. Roman et al., "Solar sintering of alumina ceramics: Microstructural development," Solar Energy, 82 (2008), 893-902. 7. D. Chakravarty et al., "Spark Plasma Sintering of Magnesia-Doped Alumina with High Hardness and Fracture Toughness," Journal of the American Ceramic Society, 91 (1) (2008), 203-208. 8. R. S. Mishra and A. K. Mukherjee, "Electric Pulse Assisted Rapid Consolidation of Ultrafine Grained Alumina Matrix Composites," Materials Science. Engineering A, A287, (2000), 178-182 9. L. Gao et al., "Bending Strength and Microstructure of Alumina Ceramics Densified by Spark Plasma Sintering," Journal of the European Ceramic Society, 20 (2000), 2149-2152. 10. S. W. Wang, L. D. Chen, and T. Hirai, "Densification of A1203 Powder Using Spark Plasma Sintering," Journal of Materials Research, 15 (4) (2000), 982-987. U.Z. Shen et al., "Spark Plasma Sintering of Alumina," Journal of the American Ceramic Society, 85 (8) (2002), 1921-1927. 12. B.N. Kim et al., "Effects of heating rate on microstructure and transparency of spark-plasma-sintered alumina," Journal of the European Ceramic Society 29 (2009), 323-327. 13. Y. Zhou et al., "Densification and grain growth in pulse electric current sintering of alumina," Journal of the European Ceramic Society, 24 (2004), 3465-3470. 14. S. J. Bennison and M. P. Harmer, "Effect of Magnesia Solute on Surface Diffusion in Sapphire and the Role of Magnesia in the Sintering of Alumina," Journal of the American Ceramic Society, 73 (4) (1990), 833-837. 15. A. H. Heuer, "The Role of MgO in the Sintering of Alumina," Journal of the American Ceramic Society, 62 (S-61) (1979), 317-318. 16. E. Sato and C. Carry, "Yttria Doping and Sintering of Submicrometer- Grained a-Alumina," Journal of the American Ceramic Society, 79 (8) (1996), 2156-2160. 17. P.Gruffel and C. Carry, "Effect of Grain Size on Yttrium Grain Boundary Segregation in Fine-Grained Alumina," Journal of the European Ceramic Society, 11 (1993), 189-199. 18. M.Stuer et al., "Transparent polycrystalline alumina using spark plasma sintering: Effect of Mg, Y and La doping," Journal of the European Ceramic Society, 30 (2010), 1335-1343.

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19. R. Voytovych et al., "The effect of yttrium on densification and grain growth in alpha-alumina," Ada Materilia, 50 (2002), 3453-3463. 20. S. Lartigue-Korinek et al., "Titanium effect on phase transformation and sintering behaviour of transition alumina," Journal of the European Ceramic Society, 26 (2006), 2219-2230. 21. H.N. Yoshimura and H. Goldenstein, "Light Scattering in polycrystalline alumina with bi-dimensionally large surface grains," Journal of the European Ceramic Society,, 29 (2009), 293-303. 22. D. Jiang et al., "Optically transparent polycrystalline AI2O3 produced by Spark Plasma Sintering," Journal of the American Ceramic Society, 91 (1)(2008), 151-154.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

SINTERING OF NANOCRYSTALLINE TUNGSTEN POWDER

William S. de Rosset U.S. Army Research Laboratory

Aberdeen Proving Ground, Maryland 21005

Keywords: nanocrystals, tungsten, powder, sintering, penetrators, modeling

Abstract

Small metal samples have been made from nanocrystalline tungsten powder by sintering with the goal of producing high density parts. The parameters associated with sintering runs have been examined to see if there is an empirical relationship between them and the final part density. A function that depends on the time the sample spends above a critical temperature has been formulated that represents one specific batch of tungsten powder. The function has been applied to other batches of powder with some success.

Introduction

The United States Army Research Laboratory (ARL) is interested in producing large tungsten parts that retain their nanocrystalline microstructure. The ultimate use for the material is in high density penetrators. It is surmised that if the parts retain their nanocrystalline microstructure, the tungsten will exhibit the same high-rate shear behavior during penetration as that of depleted uranium-3/4% titanium penetrators. This type of behavior produces superior penetration performance [1].

Nanocrystalline tungsten powder feed stock with a low (~ 1%) concentration of oxygen was purchased from Kennametal, Inc. The mean particle size was 10 microns. The powder is processed in several steps to produce a given batch. There are a large number of processing variables that go into the production of a given batch. Thus, it is expected that the properties of the compact made from each batch through cold iso-static pressing or die pressing will vary considerably. Each compact may then require a different set of sintering parameters to achieve the highest final density. Due to the sensitive nature of this information, most processing variables are not presented in this paper.

The challenge in producing a fully-dense, nano-crystalline tungsten penetrator is that the sintering process leads to a larger grain size as the compacted powder is heated. It would be advantageous if the sintering process could be modeled and the parameters governing the process could be chosen to optimize the final density while at the same time keeping the grain size as small as possible.

In the next section the considerations that went into the model development will be presented, along with fits to the data for the model. In the discussion section, the advantages and limitations of the model are brought out. An attempt is made to use the

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model to guide the choice of sintering parameters to produce the greatest density while limiting grain growth. The last section summarizes the modeling effort.

Material Specific Sintering Model

The literature was surveyed to see if viable models already existed. Most of the models found relied on a sintering time that was long in comparison to the ramp up to the sintering temperature. One such example is given in equation 1.

V = texp(-Q/kT), (1)

Here, f is the degree of sintering, / is the time held at the absolute temperature T, Q is an activation energy, and k is Boltzmann's constant [2]. It is intuitive that a longer sintering time can lead to a greater material density (i.e., degree of sintering). However, there is certainly a limit to the process. Thus, some form of time limit would be needed to apply equation 1 to a real sintering process. Also, V would have to be defined in terms of measurable parameters. A more explicit model of this form gives the sintering shrinkage KPas

r , = ^ e x p ( - ß / t o ) (2)

where Bs includes such parameters as atomic volume and vibrational frequency, D is the median particle size, and w and v depend on the diffusion mechanism [3]. Neither of these two models would apply in the case where the time to ramp up to the sintering temperature was as long as or longer than the sintering time itself.

For the present work, the first step taken in developing a simple model was to examine a representative sintering data set. The intention was to see if some analytical form could be used to fit the data. The form would be suggested by the important trends in the data. The data set chosen for the present work was all from one batch. Nineteen separate sintering trials were conducted with this batch of powder, identified as Batch A. This was the batch of powder that resulted consistently in sintered samples with small (~ 200 nm) grain sizes. A large range of parameters was used for these sintering runs, so that a better sense of the important trends could be obtained.

Results from the sintering runs made with Batch A powder are shown in table 1. The first column identifies the sintering run, and the second column shows the normalized sintering temperature, T. This normalization procedure will be discussed after presentation of the model. The third column gives the hold time, HT, at the normalized sintering temperature. The ramp rates to this temperature were such that the time to reach the final temperature was in many cases longer than the time at that temperature. Typically, the sample was cooled at twice the heating rate. The percent theoretical density was determined by weighing the sample in air and water and using 19.3 g/cm3 as the theoretical density.

The grain size is shown in the last column. Samples were prepared for viewing on the

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scanning electron microscope by polishing them with successively finer grit papers. An example of a micrograph of a sample after this final step was applied is shown in figure 1, along with an overlay grid. The grain boundary intersection count with this grid was used to determine an approximate grain size. Dividing the length of the grid by the number of intersection points gives the result. (Note: this is really the mean lineal intercept length as defined by ASTM E 112-96 (Reapproved 2004). Since there is no direct mathematical relationship between the mean lineal intercept length and the grain size number G, we simply use it as an average grain size.)

Three of the entries in table I are marked with an asterisk. These are exceptional runs that resulted in a high density (> 95% theoretical density) and a small grain size (~ 150 nm or less).

These data were examined to extract a functional relationship between the final material densities to the sintering parameters, if possible. The two major requirements on the relation were that it should incorporate all the major controllable sintering parameters (time, temperature, and heating/cooling rates) and that it have a plausible basis. The relationship should also conform to physical reality in that the final density should not exceed the theoretical density. A desirable feature of the function is that it contain as few fitting parameters as possible.

Table I. Sintering Run Data for Batch A

Sintering Identification

A-l» A-2 A-3 A-4» A-5* A-6 A-7 A-8 A-9 A-10 A- l l A-12 A-13 A-14 A-15 A-16 A-l 7 A-l 8

T, Normalized Sintering Temperature

1.628 1.512 1.512 1.628 1.628 1.628 1.628 1.628 1.512 1.628 1.512 1.628 1.295 1.744 1.744 1.512 1.744 1.860

Sintering Hold Time

1 sec 1 sec 1 sec 1 sec 1 sec 1 sec 1 sec 30 min 1 sec 4hr 4hr 30 min 30 min 30 min 1 sec 30 min 4hr 1 sec

% Theoretical Density

95.15 90.2 89.12 95.77 96.24 89.73 88.19 91.29 85.88 95.90 95.49 96.21 84.80 96.68 95.17 87.83 95.16 94.74

Grain Size (nm)

143 132 128 147 128 124 108 838 111 189 300 194 112 159 162 143 301 150

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Figure 1. Micrograph of sample A-5

An examination of the data indicated that it is important to have the sintering temperature above a certain value for a reasonable length of time in order to achieve the highest final density. Several functions were considered that fit the data to some extent and met most or all of the imposed requirements. The one chosen met all the requirements and was the simplest in form.

Let the percent theoretical density be p/p0 and t be the time the sample is above a critical temperature T0 The sintering temperature is normalized by T0 to give T. A function that describes the data over the observed range of parameters is given by

Pip0=\-\i(rl). 0)

The percent theoretical density approaches 1.0 as the values of t and x become large. Here, t is measured in minutes. Note that the time (HT) held at the maximum sintering temperature as well as both the heating and cooling rates are needed to determine t. Realistic heating and cooling rates, as well as values of r larger than 1, ensure that / is large enough to avoid negative values oîp/p„.

Data for a normalized sintering temperature of 1.628 is shown in figure 2, along with the fitted function taken from equation 3. A best fit to the data was obtained by adjusting T0. A similar plot is shown in figure 3 for r = 1.512.

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Figure 2. Comparison of data and functional fit for r = 1.628

Figure 3. Comparison of data and functional fit for r = 1.512

Discussion

The fitting function in equation 3 is restricted to one batch of powder, so its usefulness in predicting the results for sintering runs from other batches is open to question. However, the general trend of other data sets might be represented by the function. To examine this

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possibility, another batch of powder was selected that had a large number of sintering trials, Batch B. Each of the sintering runs with this batch was done at a fast heating rate, except for B-8. The runs had various hold times at the sintering temperature. Table II gives the relevant parameters for these sintering runs. None of these runs had grain sizes below 150 nm. However, the theoretical densities were all equal to or above 93% with only one exception.

The predicted densities are plotted versus the model in figure 4 for r = 1.628. The data appear to fall near the predicted curve, except for B-6. This run produced a substantially lower density sample than the identically-processed B-7. Several possible explanations were considered. However, none of them could be proven or disproven. The agreement between the predicted and observed values of the density achieved for sintering runs using Batch B is encouraging, especially considering that the fitting parameter T0 was not changed.

Table II. Sintering Parameters for Batch B

Sintering Identification

B-l B-2 B-3 B-4 B-5 B-6 B-7 B-8 B-9

T

1.512 1.628 1.628 1.744 1.860 1.628 1.628 1.628 1.628

Measured p/po

0.9470 0.9707 0.9083 0.9564 0.9624 0.9081 0.9676 0.9767 0.9679

Predicted p/p0

0.877 0.950 0.928 0.986 0.963 0.972 0.972 0.970 0.959

Grain Size (nm) 171 182 197 275 212 194 215 161 181

Sintering trials from a third batch of powder were examined, given the apparent agreement between the model and Batches A and B. The powder chosen was labeled Batch C. This batch was not processed at the ARL, so first-hand knowledge of the processing parameters was not available. However, it was believed that the parameters were close enough to those used for Batches A and B that it was worth an investigation. Initial comparisons of the data with the model were not favorable. However, upon further considerations, it was decided to use those sintering samples that were 5 g or less. This was done to be consistent with the samples used from Batches A and B, whose sample masses were generally less than 10 g. A total of six samples from Batch C were compared to the model. The results are shown in table III.

Subsequent batches were prepared with the intention of scaling the entire process up to make larger specimens. Again, for those larger sample masses the model predictions did not agree with the measured densities.

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Figure 4. Comparison of data and model for T = 1.628, Batch B

Table III. Comparison of Model Predictions with Measurements for Batch C

Sintering Identification

C-l C-2 C-3 C-4 C-5 C-6

Mass

3.26 3.39 3.19 2.34 5.03 4.75

r

1.628 1.512 1.628 1.395 1.744 1.744

Measured P/Po 0.955 0.865 0.980 0.845 0.970 0.955

Predicted P/Po

0.963 0.918 0.977 0.788 0.982 0.980

Difference (%) 0.8 5.8 0.3 7.3 1.2 2.6

While no definitive statement can be made concerning the maximum sintering temperature, it can be brought into the analysis in the following way. First, we select a desired value ofp/p0. For our example, 0.95 is chosen. This is an arbitrary choice, and the analysis can be carried out for other values. In this case,

0.95 = l - l / ( r r" ' ) (4)

from equation 3. Assuming a cooling rate that is twice the heating rate, values of t and x can be generated for given values of HT and HR '. We arbitrarily set the value

of HT to 30 min and examine the relationship among t, t, and HR '. These results are shown in table IV. In addition to minimizing t, it is reasonable to assume that small grain growth will occur for a minimum value of T. However, according to the model and our assumptions concerning the other parameters, t and T cannot be minimized simultaneously.

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Table IV. Values off and r from Model Predictions

HR-(min-1)

20 22.5 25 27.5 30 32.5 35

t (min)

74.8 70.4 66.8 63.8 61.3 59.1 57.3

T

1.694 1.705 1.713 1.721 1.728 1.734 1.740

Summary

An empirical approach has been used to organize sintering data for a specific batch of powder. One functional form that appeared to fit the data for s specific batch of powder was based on the time t the powder sample was held above a critical temperature, Ta, used as a single fitting parameter. The same function and fitting parameter were used to fit the data from another batch with reasonable success. Good agreement between the model predictions and measured densities was found with a third batch of powder, so long as the mass of the samples was kept at 5 g and below. This work has shown that if attention is paid to keeping the powder processing parameters the same as well as the sample masses, the possibility exists that the final density can be predicted using values of the controllable sintering parameters, based on a model for one single batch of powder.

Acknowledgments

The author is indebted to Eric Klier and Brady Butler for many helpful discussions on sintering theory and data interpretation. They were the ones responsible for generating the data found in this report. Thanks also go to James Catalano for providing a final polishing step that brought out the grain structure in the micrographs of the nanocrystalline tungsten and made the grain size determination more accurate. Finally, the author is grateful for the micrographs taken by Bradley Klotz.

References

1. Lee S. Magness, "High Strain Rate Deformation Behaviors of Kinetic Energy Penetrator Materials During Ballistic Impact," Mechanics of Materials, 17 (1994), 147-154. 2. Randall M. German, Sintering Theory and Practice, (John Wiley and Sons, New York, 1996.) 3. R.M. German, and E. Olevsky, "Strength predictions for bulk structures fabricated from nanoscale tungsten powders," International Journal of Refractory Metals and Hard Materials, 23 (2005), 77-84.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF POWDER SYNTHESIS AND PROCESSING ON LUMINESCENCE PROPERTIES

J. McKittrick1'2, J.K. Han2, J.I. Choi2 and JB. Talbot2'3

University of California, San Diego, La Jolla, CA 92037 USA 'Department of Mechanical and Aerospace Engineering

2Materials Science and Engineering Program 3Nanoengineering Department

Keywords: powder synthesis, powder processing, luminescent properties

Abstract

Visible light-emitting materials (phosphors) have broad use in a variety of applications, from fluorescent and LED lighting, displays, and scintillators. Powder synthesis and post-processing methods significantly affect the luminescence properties. Several methods are outlined and compared. The luminescence properties can be used to evaluate the quality of powders - the chemical composition and crystallite size.

Introduction

Phosphors are luminescent materials that emit photons (IR to UV) when stimulated by an external energy source (photons, electrons, electric fields, etc.). Phosphors have wide applicability for lighting, in display screens, as x-ray counters (scintillators), safety stripes, and more recently for solid state lighting. Phosphors used for their photoluminescence properties (e.g. fluorescent and LED lighting) are typically wide band gap oxides or nitrides whereas sulfide-based semiconductor compositions have applicability for electroluminescence devices. These phosphors have two components - an inert host lattice and a small quantity of an activator, which is a transition or rare earth element. The activator does all the 'work' in producing the emissive visible-range photons (400-700 nm) - the oxide lattice is simply a cage that keeps the activators separate from each other and provides a stable environment. Compositions are written as host:activator. There are hundreds of phosphor compositions that have been optimized for specific applications (see for example the Phosphor Handbook [1]).

One way to evaluate the physical and chemical properties of oxide ceramic powders is through measurement of the photoluminescence emission intensity and wavelength, which can be used to identify the crystallite size and composition, respectively. Along with a discussion of applications of luminescence materials, this report also describes powder preparation methods and the resultant effect on the luminescence properties. A focus is placed on oxide hosts for use in photoluminescence applications.

Lighting applications and phosphors

In the case of fluorescent lighting and plasma display panels (PDP), the phosphors (which are coated on the inside of the fluorescent tube or on a dielectric for PDPs) are activated by 254 nm (4.9 eV) photons from the Hg discharge in the plasma in the lamp or pixel (for PDPs). Phosphors development for this application is a mature field with few new compositions discovered each year. For example, the blue-emitting phosphor in PDPs is BaMg2Ali6027

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activated by Eu +. To produce white light, a single composition that emits a broad band in the visible or a tri-blend of red-, green- and blue-emitting compositions can be used for fluorescent lamps. For PDPs, red-, green- and blue-emitting stripes are situated side-by-side in each pixel to give a full color display.

For solid state lighting (LED lamp), an InGaN blue-emitting diode is used, which is coated with a yellow-emitting phosphor (Y3Al50i2:Ce3+) that in combination produces white light. However, the white 'color' is not optimal, since it lacks good emission in the red portion of the spectrum, as shown in Figure 1(a).

Figure 1. (a) CIE color coordinate diagram, showing that a mixture of blue and yellow yields white, (b) spectral energy distribution of a blue-emitting LED chip in combination with a yellow-emitting phosphor (Y3Al5Oi2:Ce3+) results in emission across the visible spectrum (400-700 nm) and (c) optical micrograph of a cross-section of a commercial blue-emitting LED with phosphor particles surrounding it. (Scale bar = 50 fun)

Figure 2. Relationship between the crystallite size and the relative emission intensity of Y2Û3:Eu phosphor particles prepared by spray pyrolysis. Adapted from [2],

The most important properties of phosphors are the photoluminescence emission and powder characteristics. The quantum efficiency (n = # photons in/# photons out), emissive energy distribution (plot of emission intensity as a function of wavelength (Figure 1(b)) and thermal stability (slight changes in temperature can severely affect r\) chiefly factor in the selection of a particular composition. The particle size, morphology and degree of agglomeration are the most important powder properties. Figure 1(c) shows an optical micrograph of a commercially available white-emitting LED illustrating the non-uniform particle dispersion and large particle size distribution. The powders are mixed in a polymer to create a dome that is placed over the LED chip. Most phosphors are used as powders,

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although there are applications for thin films and single crystals, notably in electroluminescent displays and as scintillators, respectively. The particle size and morphology are important since most phosphors are not used as dispersed powders (as in white-emitting LEDs), but are deposited on a substrate (by settling or electrophoretic deposition). Smaller particles result in the use of less material, a higher resolution screen and a reduction of scattering. Thus, there is great interest in developing nanocrystalline phosphors.

Figure 3. Photoluminescence emission spectra as a function of wavelength of core and core (Y203:Eu3+) / shell (SiC>2) nanophosphors for different shell thicknesses. Adapted from [6].

The problem with nanocrystalline phosphors is that the emission intensity (correlated to the quantum efficiency) increases as the crystallite size increases (Figure 2). This has been attributed to surface defect surface states that trap the emitted photons and thus quench the emission intensity. These defects arise from dangling bonds, non-uniform crystalline environment around the activator, and absorbed species. Another possibility is lattice distortion that occurs with small crystallite sizes due to the poor crystallization. Although it is relatively easy to prepare nanocrystalline powders by use of liquid precursors, they must be calcined to remove water and organics, which results in sub-micron sized, sintered agglomerates that are not easily separated into individual crystallites. The nanocrystallinity of the as-synthesized powders make them relatively easy to sinter at low temperatures (for a ceramic) and the presence of the small amount of activator may accelerate the sintering rate [3,4].

Figure 4. SEM micrographs of the core/shell particles for different SiC>2 deposition times: (a) 2 hr (b) 4 hr (Scale bar=100nm). (c) Laser particle size distribution. Adapted from [6].

One method ameliorate surface defects is to coat the nanoparticles with an inert shell [5,6]. Figure 3 shows that the emission intensity can be significantly improved by coating Si02 onto Y203:Eu nanophosphors. Figure 4(a) and (b) show SEM micrographs of the core/shell nanophosphors - the bright cores (high atomic number) are surrounded by the darker shells. Also of significance is that only one or two particles are coated thereby producing isolated cores. The uniformity of the shell thickness results in spherical, uniformly shaped particles with a narrow particle size distribution, as shown in Figure 4(c).

The composition has an effect on the powder morphology. Figure 5(a-d) shows SEM micrographs of (Sri.xBax)2Si04:Eu2+. As the

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x increases, the particle size decreases. Compositional uniformity is very important for phosphors - Figure 5(e) shows how the peak emission wavelength changes by varying x. Thus using the emission spectra, the composition of powders can be accurately determined.

Figure 5. SEM micrographs of (Ba^Sr^SiCvEu of for (a) x = 0, (b) x = 0.25, (c) x = 0.75 and (d) x = 1. Scale bar = 500 run. (e) Photoluminescence emission spectra of the powders.

Powder synthesis methods

Many ceramic powder synthesis and processing techniques have been applied to phosphor development - the most popular methods include sol gel, co-precipitation, hydrothermal, spray pyrolysis, solid state reaction, combustion synthesis and the fabrication of combinatorial libraries. Most phosphor compositions have > 4 elements, which make fabrication of uniform, single phase powders difficult. Selection of a synthesis method depends on the goal. If low cost preparation is the goal, solid state reactions are favored. This involves heating precursor oxides or nitrides at high temperatures and allowing interdiffuion to occur. In most cases, due to the multiple precursors, several sintering and grinding steps are required to get a full reaction. The resultant body is then ground or milled to micron-size (3-10 urn) powders. Extensive grinding of phosphors needs to be avoided because the induced lattice defects reduce the quantum efficiency. If small, uniform shape, narrow particle size distribution, dispersed powders are the goal, then liquid or gaseous methods should be employed. The drawback to these methods is the cost, lengthy synthesis time and the small volume of solid produced. If evaluation of a large number of compositions is the goal, then combinatorial synthesis or combustion reactions are preferred.

Combustion synthesis is a popular method to produced multiconstituent oxides due to the fast preparation time. It involves igniting an aqueous solution of metal nitrates and a fuel at a low temperature (< 50CfC) and allowing the highly exothermic, oxidative reaction to occur, usually with an accompanying flame. This results in reaction temperatures of > 1200 "C and reaction times < 10 min [7-9], The reaction proceeds by first boiling off the water that produces a gelatinous mass, which then bursts into flame. Shown below is the reaction to form a red-emitting phosphor, (Y0.95EU0 05)203 using nitrate precursors and a fuel (urea):

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1.9Y(N03)3 + 0.1Eu(NO3)3 + 5CO(NH2)2 -> (Yo.9sEuo.o5)203 + 10H2O + 8N2 + 5CO:

Due to the large amount of gas produced, the product is a highly porous, crystalline mass that is easily reduced to powder form by light crushing. The flame temperature can be adjusted by using different fuels. Figure 6(a) shows the photoluminescence emission intensity dependence on the measured flame temperature of red-emitting Y3(Ali.xCrx)sOi2. Thermodynamic analysis of the reaction demonstrated that carbohydrazide fuel was the most exothermic, thus produced the highest flame temperature and most luminescent powders [9]. Due to the short duration at the high temperature, nanocrystalline powders are produced (measured by TEM analysis and x-ray line broadening), however they are agglomerated into hard, sub-micron sized powders that are not easily separated (Figure 7). The crystallite size is larger at the higher reaction temperatures that has promoted of crystallite growth. Figure 6(b) shows the effect of post-synthesis thermal treatments on the emission intensity. Although the flame temperature is > 1800'C (carbohydrazide fuel) the emission intensity increases with further heat treatments. This is attributed to two factors - removal of excess entrapped carbon and further crystallite growth.

Figure 6. (a) Photoluminescence emission spectra and measured flame temperatures of Y3(Ali. iCrx)50i2 phosphors made with either urea, carbohydrazide or glycine fuels. Taken from [8]. (b) Effect of post-synthesis heat treatments on the photoluminescence emission intensity of YAG:Cr produced with carbohydrazide fuel. Taken from [9], Matijevic was one of the first to observe that powder morphology could be widely varied depending on the chemical precursors [11]. An illustration is shown in Figure 8, where needle-like or spherical particles of Sr2Si2C>4:Eu2+ were achieved by co-precipitation, combustion or sol gel methods. For phosphor applications, it is desirable to produce uniform, spherical phosphors, thus the selection of a synthesis method is the first crucial step in phosphor development. A comparison of the various synthesis methods used to prepare the powders in Figure 8 is outlined below.

In co-precipitation, salts of the precursor oxides are added along with tetraethyl orthosilicate (TEOS) are added to dimethylformamide and stirred. Sodium hydroxide dissolved in a water/ethanol mixture is added, resulting in the nucleation of small particles in solution, which are then filtered, dried and heated between > 600°C to crystallize.

A modified sol gel/Pechini method is as follows: metal oxides in the appropriate molar ratios are first dissolved in a nitric acid solution (cheaper than buying metal nitrate powders). Citric acid, a

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chealating agent and ethylene glycol are then added along with polyethylene glycol (cross linking agent). Stirring at a temperature < lOCfC results in a transparent gel, which is then calcined to remove organics and water, resulting in nanocrystalline powders. These particles have a uniform, narrow particle size distribution, as opposed to combustion synthesis. However, the emission intensity is strongly dependent on the post-synthesis annealing temperature, as shown in Figure 9, which is attributed to the increase in crystallite size. There is an abrupt increase in emission intensity and particle size at temperatures from 1100-1200'C, indicating that the reduction in surface area is largely responsible for the high emission intensity.

Figure 7. Dark field TEM micrographs of combustion synthesized blue-emitting nanocrystalline Y2Si05iCe3+ showing influence of crystal structure and morphology of the hard agglomerates: (a) low temperature XI phase and (b) high temperature X2 phase. Adapted from [10].

The calcination step (usually between 800-1200°C) required for all these synthesis methods results in particle sintering (neck formation at lower temperatures, boundary sintering at higher temperatures) and is accompanied by crystallite growth. Thus there is a tradeoff between obtaining dispersed, uniform size and shape nanocrystalline phosphors and maximizing the emission intensity. For all commercial applications, emission intensity cannot be sacrificed. It remains an open challenge to fabricate dispersed nanocrystalline phosphors with high quantum efficiency.

Figure 8. SEM micrographs of (Sro.97Euo.o3)2Si2C>4 green/yellow-emitting particles prepared by (a) co-precipitation with dimethylformamide solvent, (b) combustion synthesis with urea fuel and (c) sol gel method. (Scale bar = 500 nm)

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Figure 9. Integrated photoluminescence emission intensity as a function of post-synthesis annealing temperature on sol gel prepared powders of green-emitting Ba2SiC>4:Eu . The quantum efficiency = 95% at Normalized Emission Intensity = 1.

Conclusions

Luminescent materials (phosphors) development require careful selection of the synthesis method and post-synthesis processing treatments. The most important phosphor properties are the intensity and shape of the spectral energy distribution and the powder characteristics. Nanocrysalline oxide phosphors are desired for many applications, however the quantum efficiency is low and a calcination step is required to improve the luminous output. This step sinters the particles into agglomerates that are hard to separate into individual nanocrystals. Luminescence properties can be used to evaluate the quality of powders, including the composition and crystallite size.

Acknowledgments

We thank Drs. Kailash Mishra, Mark Hannah, Alan Piquette and Maria Anc from OSRAM/Sylvania, Central Research, Beverly, MA for helpful discussions. This work is supported by the Dept. of Energy, Grant DE-EE-0002003.

References

1. S. Shionoya, W.M. Yen and H. Yamamoto (Eds.), Phosphor Handbook, 2nd Edition (Boca Raton, FL: CRC Press, 2006).

2. K.Y. Jung, C.H. Lee, Y.C. Kang, "Effect of surface area and crystallite size on luminescent intensity of Y2Û3:Eu phosphor prepared by spray pyrolysis," Mater. Lett., 59 (2005) 2451-2456.

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3. R.M. German, Powder Injection Molding (Princeton, NJ: Metal Powder Industries Federation, 1990).

4. R.M. German and Z.A. Munir, "Enhanced low-temperature sintering of tungsten," Metall. & Mater. Trans. A, 7 (1976) 1873-1877.

5. K. Kompe, H. Borchert, J. Storz, A. Lobo, S. Adam, T. Möller and M. Haase, "Green-emitting CeP04:Tb/LaP04 core-shell nanoparticles with 70% photoluminescence quantum yield," Angew. Chem. Int. Ed., 42 (2003) 5513-5516.

6. J.K. Han, G.A. Hirata, J.B. Talbot and J. McKittrick, "Luminescence enhancement of Y203:Eu3+ and Y2Si05:Ce3+,Tb3+ core particles with Si02 shells," Mater. Sei. & Eng. B, 176(2011)436-441.

7. J.J. Kingsley and K.C. Patil, "A novel combustion process for the synthesis of fine particle a-alumina and related oxide materials," Mater. Lett., 6 (1988) 427-432.

8. J. McKittrick, L.E. Shea, CF. Bacalski and E.J. Bosze, "The influence of processing parameters on luminescent oxides produced by combustion synthesis," Displays, 19 (1999) 169-172.

9. L.E. Shea, J. McKittrick and O.A. Lopez, "Synthesis of red-emitting, small particle size luminescent oxides using an optimized combustion process," J. Am. Ceram. Soc, 79 (1996)3257-65.

10. E.J. Bosze, J. McKittrick and G.A. Hirata, "Investigation of the physical properties of a blue-emitting phosphor produced using a rapid exothermic reaction," Mater. Sei. Eng. B, 97 (2003) 265-273.

11. E. Matijevic and P. Scheiner, "Ferric hydrous oxide sols II. 1. Preparation of uniform particles by hydrolysis of Fe(lII)-chloride, -nitrate and -perchlorate solutions, J. Coll. & Interface Sei., 63 (1978) 509-524.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF RAPID SOLIDIFICATION AND HEAT TREATMENT ON D2 TOOL STEEL

Pooya Delshad Khatibi, Douglas G. Ivey, Hani Henein

Chemical and Materials Engineering Department, University of Alberta; Edmonton, Alberta, T6G 2V4, Canada

Keywords: D2 tool steel, rapid solidification, M7C3 carbides

Abstract Improved mechanical properties for D2 steels can be achieved by reducing microsegregation and obtaining a good distribution of carbides during solidification. An understanding of the evolution of microsegregation and carbide formation during rapid solidification is necessary to control the microstructure. Impulse atomization (IA) was employed to produce rapidly solidified (RS) powders of D2 steel in helium and nitrogen cooling gases. The resulting powders were sieved into different size ranges and the microstructures and phases were examined using SEM and XRD. In addition to a substantial refinement of grain size, rapid solidification also suppressed the martensitic transformation and produced supersaturated retained-austenite. Several annealing tests were carried out on the as-atomized powders. SEM images showed a good distribution of precipitated carbides after annealing. The effect of annealing on the RS powders was evaluated using Vickers microhardness measurements and the precipitated carbides were characterized using TEM.

Introduction One of the well-known and widely used tool steels in industry is D2 tool steel. D2 tool steel has a high volume fraction of carbides, which results in a good wear and abrasion properties. In order to obtain superior mechanical properties, microsegregation should be minimized [1]. Conventional casting of tool steels will result in a coarse carbide structure. High chromium, high carbon tool steels contain a hard (Fe,Cr)7C3 phase with a hardness in the 1300-1800 HV range. The abrasion resistance of these alloys is not very good, however, because the carbides formed under conventional solidification conditions are coarse. Therefore, a rapid solidification technique is used for high chromium, high carbon tool steels to refine the microstructure [2], Blaha et al showed that cold worked tool steel, produced via a powder metallurgy process, possessed improved mechanical properties due to the fine distribution of carbides [3].

Although D2 steel is one of the most popular and commercially important die steels, so far microstructural characterization has only been reported for RS material using splat quenching and melt spinning [4,5]. In the current study, the effect of rapid solidification via impulse atomization (IA) and subsequent heat treatment on the microstructure and phase formation of D2 tool steel is analyzed. Based on XRD and TEM studies, it is shown that RS results in the formation of a supersaturated metastable austenite phase instead of the equilibrium ferrite phase. Also, annealing of atomized particles results in formation of fine rod-like carbide precipitates inside a ferrite matrix.

Experimental Table I shows the chemical composition of the D2 tool steel, which was used in this study.

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Table I. D2 tool steel nominal chemical composition (wt.%)

c 1.55

Cr

11.8

Mn

0.40

Mo

0.80

Si

0.30

V

0.80

Fe

Balance

In order to understand the different phases that form during solidification of a D2 tool steel, the phase diagram for this steel was calculated by means of ThermoCalc software and the TCFE6 data bank (Fig. 1). The arrow in this figure indicates the chemical composition of the D2 tool steel evaluated in this paper.

Figure 1. Calculated phase diagram for D2 tool steel.

In order to rapidly solidify the D2 tool steel, an impulse atomization (IA) process, as explained elsewhere [6] was used. Impulse atomization was performed by holding the D2 tool steel at 1600°C for about 30 minutes. Then, the liquid was pushed through orifices at the bottom of the crucible and metal droplets were generated. Falling droplets were cooled in helium and nitrogen atmospheres with a maximum oxygen content of 10 ppm. The powders were solidified by the time they reached an oil quench bath 4 m below the atomizing nozzle. Solidified particles were washed using toluene and methanol and then sieved into different size ranges based on MPIF Standard 05 [7]. Particles in the 600-710 um size range, solidified in both helium and nitrogen atmospheres, were chosen for further analysis (hardness measurements, XRD, SEM and TEM).

To demonstrate that the microstructure in the as-atomized powders was unstable, the helium atomized particles (600-710 um) were given an isochronal heat treatment for 1 hour at temperatures from 623 to 1083 K and subsequently cooled to the room temperature. Annealing was done in a Setaram Labsys Evo Differential Scanning Calorimeter.

Samples of chosen particle sizes were mounted in epoxy and then ground and polished for microstructural analysis. Samples were then carbon coated in preparation for microstructural analysis using scanning electron microscopy (SEM). A Zeiss Evo MA 15 SEM with a 20 keV electron beam was used for microstructural analysis. EDX analysis in the SEM was done using a

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Bruker Silicon Drifted Detector with an energy resolution of 125 eV. Phase identification was carried out through X-ray diffraction using Co-Ka radiation in a Rigaku Denki Rotâlex RU -200B X-ray system and neutron diffraction (ND) using a neutron beam of 1.33A" wavelength. ND was performed at Atomic Energy of Canada Limited (AECL) in Chalk River. ND was used to measure the austenite lattice parameter of as-atomized particles (600-710 um); lattice parameters were extracted using GSAS software. As-atomized powders (in helium) and atomized powders annealed at 923 and 1023 K were also characterized by transmission electron microscopy (JEOL 2010 TEM operated at 200 kV). Electron transparent specimens were prepared using focused ion beam (FIB) techniques (Hitachi NB 5000, dual-beam FIB/SEM).

Hardness measurements were also done on the as-atomized and annealed particles using a MVK-Hl Mitutoyo Vickers hardness tester.

Results and Discussion Fig. 2 shows the SEM microstructures for two IA powders of D2 tool steel in the 600-710 urn size range solidified in helium (Fig. 2a) and nitrogen (Fig. 2b). Light and dark gray phases as well as some solidification porosity are visible in both images. From previous work, the light gray phase was likely primary retained-austenite and the dark region was likely the eutectic composed of carbide and austenite [8]. Visual observation of Fig. 2a and 2b shows that the microstructure is finer for the particles solidified in helium versus those solidified in nitrogen, which is because of the higher thermal conductivity for helium compared with nitrogen [9], It was shown previously that the cooling rate for IA particles in helium and nitrogen varies from 102 to 104 K/s [8].

Figure 2. SEM backscattered electron (BSE) images from two particles of D2 tool steel: (a) Atomized in helium and (b) atomized in nitrogen. The particle size was 655 um for both samples.

Fig. 3 shows the XRD results for as-atomized particles (600-710 um) and heat treated particles at different temperatures. It can be seen that rapid solidification of D2 steel causes formation of supersaturated metastable retained-austenite and complete suppression of ferrite and martensite formation. The austenite in the rapidly solidified D2 steel particles was so stable that even after quenching in liquid nitrogen, no austenite was converted to martensite. The reason is likely the lowering of the martensite start temperature caused by high supersaturation of austenite. Bhargava et al showed that the lattice parameter of austenite in D2 steel, rapidly solidified by the chill block melt spinning method, is about 1.5% larger than that for y-Fe, indicating substantial dissolution of carbon and chromium in the austenite [4]. Annealing the particles at 623 K results in no change to the phases, but upon annealing at 923 K retained austenite has transformed to

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ferrite. In addition, low intensity peaks, which can be tentatively indexed to M7C3, are present. The austenite to ferrite transformation, then, has resulted in carbon precipitation.

Figure 3. XRD patterns for 600-710 urn particles atomized in helium and annealed for 1 hour at different temperatures.

The effect of isochronal (1 hour) heat treatment on the microhardness of IA D2 steel is shown in Fig. 4. There is a hardening peak at about 923 K, which is due to significant carbide precipitation at this temperature. The hardness plots show an initial decrease in hardness, likely due to relief of internal stresses introduced by rapid solidification. Annealing at temperatures higher than 923 K. also results in a decrease in the hardness due to coarsening of carbide precipitates [4].

Figure 4. Vickers hardness for powders atomized in helium and nitrogen and annealed for 1 hour at different temperatures.

Fig. 5a and 5b show the microstructure of 600-710 urn particles annealed at 923 and 1023 K, respectively. The microstructure of the particle annealed at 1023 K shows some rod-shape precipitates inside the primary phase. In the particle annealed at 923 K, upon closer examination, there are some features within the primary phase which are barely visible. According to Fig. 4, the maximum hardness was achieved by annealing at 923 K, so precipitation is likely to have

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occurred at this temperature. Higher magnification analysis was needed to characterize the microstructure of the 923 K annealed particle.

Figure 5. SEM BSE images of D2 tool steel, atomized in helium: (a) Annealed for 1 hour at 923 K and (b) annealed for 1 hour at 1023 K. The particle size is 655 urn.

More detailed microstructural analysis was done using TEM of FIB specimens extracted from specific regions in the various microstructures. Fig. 6 shows a TEM bright field (BF) micrograph from the interdendritic region of a particle annealed at 623 K (Fig. 6a) and a selected area diffraction (SAD) pattern from the matrix or primary phase (Fig. 6b). The diffraction pattern can be indexed to austenite with an orientation close to the [112] zone axis. The austenite primarily consists of Fe with about 10 wt.% Cr and smaller amounts of V and Si (<1 wt.% for both). No precipitation has occurred in the primary austenite phase at this annealing temperature.

Figure 6. (a) TEM BF image of the interdendritic region of a 655 um particle of D2 tool steel, atomized in helium and annealed for 1 hour at 623 K. (b) SAD pattern from the austenite primary phase.

The interdendritic region represents the eutectic component, which consists of two phases, one being austenite. The interdendritic (eutectic) region is shown at higher magnification and a different tilt angle in Fig. 7. The second phase is Fe and Cr rich (~33 wt.% Cr) with smaller but

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significant amounts of V and Mo (4-5 wt.% of both). The SAD pattern in Fig. 7b is from both the austenite and the second phase, so there are two overlapping patterns. The spots corresponding to austenite are circled in red (the orientation is close to a [013] zone axis). The other spots are from the second phase, which can be indexed to M7C3 (close to a [1 j.7] zone axis), where M is Fe and Cr, as well as V and Mo. These two phases, austenite and M7C3, are consistent with previous work [10].

Figure 7. (a) Higher magnification TEM BF image from the eutectic area of the sample annealed at 623 K (Fig. 6); (b) SAD pattern from the interdendritic (eutectic) region.

Fig. 8 shows a TEM BF micrograph of a 600-710 urn particle annealed at 923 K. An SEM image from the same sample was also shown in Fig. 5a. The original primary austenite phase, which has transformed to ferrite and carbide precipitates (after annealing), and the eutectic region are indicated by arrows in Fig. 8. A higher magnification image of the transformed phase is shown in Fig. 8b. After annealing at 923 K, rod-like carbides have precipitated wimin trie primary phase; these were not visible in the SEM image (Fig. 5a). In fact, the primary phase is no longer austenite, but ferrite plus the precipitates. The precipitates are 200-300 run long and less than 50 nm in width. SAD patterns from a ferrite grain and the carbides precipitated within the ferrite grain are shown in Fig. 8c and 8d, respectively. Fig. 8d shows a partial ring pattern for M7C3 carbides as well as spots from the ferrite matrix. The carbides are Fe and Cr rich (with about 24 wt.% Cr) and smaller amounts of Mo and V (-0.4 wt.% for both) in substitution^ solid solution) and the ring pattern can be indexed to M7C3. Patterns were also obtained from other regions and at other orientations and confirmed the identity of the carbides. Although SAD patterns show that the carbides in as-atomized particles (interdendritic area) and particles annealed at 923 K (precipitated during annealing) are both M7C3, EDX results show that interdendritic carbides in the as-atomized sample have higher amounts of Cr (-33 wt.% Cr) compared with carbides precipitated within the transformed austenite grains (-24 wt.% Cr). EDX results also show that the amount of V dissolved in the matrix phase decreases from -0.6 wt.% in the as-atomized sample to -0.2 wt.% in the annealed sample. The excess V is present in the precipitated carbides.

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Figure 8. (a) and (b) TEM-BF images of a 655um D2 tool steel sample, annealed at 923 K. SAD patterns from (c) a ferrite grain; (d) ferrite and carbide.

A large lattice parameter (0.3618 nm) was measured from the ND patterns, for metastable austenite in rapidly solidified IA particles. Comparison of this value with the 0.3572 nm value for retained austenite in conventionally processed and hardened D2 steel [4] suggests that, even at cooling rates of 103 K/s, there is a high degree of supersaturation of austenite with respect to solute content, especially C and Cr. Bhargava [5] also showed that in the melt-spun method, clustering of interstitial solute atoms happens in the vicinity of dislocations within the dendritic region.

Conclusions Rapid solidification was applied to a D2 tool steel using impulse atomization (IA). It was shown that even at cooling rates of 103 K/s, metastable austenite is the primary phase containing supersaturated solute elements. Annealing at different temperatures was done on the as-atomized particles. Microhardness results showed that annealing resulted in an increase in hardness for rapidly solidified D2 steel particles. XRD analysis showed that during annealing primary retained austenite, supersaturated in alloying elements and carbon, transformed to ferrite and carbide precipitates which were responsible for the higher hardness. TEM analysis confirmed

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that as-atomized particles and those annealed at low temperature (623 K.) consisted of austenite primary phase and austenite plus M7C3 in the eutectic. Annealing at higher temperatures (>923 K) resulted in decomposition of me austenite primary phase to ferrite plus M7C3. Austenite in the eutectic area also transferred to ferrite at 923 K. Fine rod-shaped M7C3 carbides, precipitated in the matrix, were responsible for the increase in hardness.

Acknowledgments The authors would like to thank the Natural Sciences and Engineering Research Council of Canada (NSERC) and the Canadian Space Agency (CSA) for the financial support. The authors are also grateful to M. Kupsta and the National Research Council - National Institute for Nanotechnology for preparing the FIB samples.

References 1. G. Roberts, G. Krauss and R. Kennedy, Tool Steels, (5th Edition, ASM International Publications, 2000). 2. Y. Dai et al., "Solidification Structure of C2.osCr25«Si 1 j9Mno43Fe70.87 Powders Fabricated by High Pressure Gas Atomization", Materials Characterization, 61(1) (2010), 116-122. 3. J. Blaha, C. Krempaszky and E.A. Werne, "Carbide Distribution Effects in Cold Work Tool Steels" (Paper presented at 6* International Tooling Conference, Karlstad, 2002). 4. A.K. Bhargava and A.N. Tiwari, "Effect of Rapid Solidification and Heat Treatment on D2 Steel", Internationaljournal of Rapid Solidification, 7 (1992), 51-66. 5. A.K. Bhargava and A.N. Tiwari, "Some Microstructure Aspects of Melt-Spun D2 Steel", Transactions of the Indian Institute of Metals, 58 (2005), no. 1:41-47. 6. H. Henein, "Single Fluid Atomization through the Application of Impulses to a Melt", Materials Science and Engineering, 326 (2002), 92-100. 7. "Standard Test Methods for Metal Powders and Powder Metallurgy Products", Metal Powder Industries Federation, (1993), Princeton, NJ. 8. P. Delshad Khatibi, H. Henein and A Ilbagi, "Microstructural Investigation of D2 Tool Steel during Rapid Solidification Using Impulse Atomization" (Paper presented at TMS 2011 Conference, San Diego, California, 2011). 9. W.M. Rohsenow and J.P. Hartnett, Handbook of Heat Transfer, (3rd Edition, McGraw-Hill, 1998). 10. X. Wu and G. Chen, "Microstructural Characteristics and Carbide Transformation of Laser-Cladded Fe-Cr-W-Ni-C Coatings during High-Temperature Tempering", Journal of Materials Science Letters, 17 (1998), 1849-1852.

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DEVELOPMENT OF SOLID FREEFORM FABRICATION FOR METALLIC PARTS USING SELECTIVE INHIBITION OF SINTERING

M. Yoozbashizadeh, B. Khoshnevis

Industrial & Systems Engineering Department University of Southern California, Los Angeles, CA 90089, USA

Keywords: Metal parts fabrication, rapid manufacturing, powder sintering and selective inhibition of sintering

Abstract

The SIS-Metal process has been developed based on the microscopic mechanical inhibition principle. In this process metal salt (such as aluminum sulfate or Aluminum Nitrate) or water-soluble crystalline carbohydrates (such as sugar) is printed in the selected areas of each metal powder( Bronze and Titanium) layer; salt crystals decompose to metal oxide and; during the sintering process dispersed particles of metal oxides that spread uniformly between metal powder particles prevent the fusing of these particles together; consequently, the sintering process in the affected regions is inhibited. Experimental results are also presented to demonstrate the capability of the process in fabricating metal parts with various geometries and for different metals such as Titanium.

Introduction

The Selective Inhibition Sintering (SIS) process is an additive manufacturing (AM) technology which builds parts on a layer-by-layer basis. The principle idea of the SIS process is the prevention of selected segments of each powder layer from sintering. Therefore, the SIS process may be considered as an opposite approach to the Selective Laser Sintering (SLS) process in which selected areas of powder are sintered by a fine laser beam. The inhibition mechanism used in the SIS process plays the major role in successfully developing the process. There are four possible inhibition mechanisms in SIS: (1) macroscopic mechanical inhibition; (2) microscopic mechanical inhibition; (3) chemical inhibition; and (4) thermal inhibition. In this paper the microscopic mechanical inhibition for metallic parts will be investigated.

SIS-Metal Process Based On Microscopic Mechanical Inhibition

The SIS-Metal process based on microscopic mechanical inhibition is shown in Figure 1 .A bulk sintering approach is adopted for SIS-Metal. In addition, the layer-by-layer sintering of metal powder would require the fabrication machine to operate in an oxygen-free environment. Undesirable shrinkage effect among sintered layers may be generated in the layer-by-layer sintering approach, which would create undesired internal stresses.

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Figure 1. The SIS-Metal process based on the microscopic mechanical inhibition

The SIS-Metal process based on the bulk sintering approach has five main steps as follows: (1) Printing sintering inhibitor: A deposition nozzle with a fine orifice, or an inkjet print head,

is used to deliver a sintering inhibitor to the selected areas (2) Laying a thin powder layer. Metal powder is spread as a thin layer over the build tank using a

blade or a roller. (3) Creation of a boundary to contain part: A consolidating liquid is deposited on the powder

bed at the periphery of the part profile (4) Heating: A heater is used to evaporate water and other liquid additives (such as alcohol

which may be needed for breaking surface tension) in the printed inhibition sections. (5) Bulk sintering in an oven: After all the layers have been completed, a metal powder block is

extracted from the build tank and placed in a conventional sintering oven. After sintering and cooling, the sintered metal block is removed from the oven. Finally, the fabricated part can be extracted by removing the regions that are isolated by the inhibited powder areas.

Advantages Of SIS-Metal:

Metallic parts fabrication based on metal casting and powder metallurgy (P/M) requires tooling construction [7]. However, the construction of such tooling can be expensive and time consuming, especially for small quantities of parts or parts with complex geometry. In comparison, additive manufacturing (AM) processes can be much faster and less expensive for such small lot or complex part fabrication.

Currently, many AM processes have been developed for building metallic parts, such as Selective Laser Sintering (SLS), 3D Printing (3DP), Fused Deposition Modeling (FDM), Direct Metal Laser Sintering (DMLS), Laser Engineered Net Shaping (LENS), and Electric Beam Melting (EBM). These processes can be classified based on the usage of binders and the sintering approach as shown in Figure 2.

Based on such a classification the SIS-Metal process is uniquely positioned among all the metallic part fabrication processes because it is a layered fabrication process which is based on the bulk sintering approach without using any binder. The SIS-Metal process has the following advantages. • The hardware of the SIS-Metal process can be inexpensive. • The SIS-Metal process is potentially fast.

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The SIS-Metal process can have less shrinkage and deformation since no de-binding is involved. Unlike LENS or DMLS, no complex supports are required in the SIS-Metal process since overhang features are supported by powder volumes underneath.

Figure 2. A classification of the metallic part fabrication processes

Experimental Procedure

A three-axis prototype machine has been designed developed as shown in Figure 3. In the machine a single print nozzle is used for printing the inhibitor. The nozzle has a continuous speed of 31 mm/second. A heater has been incorporated for heating up every layer after printing the inhibitor. The heater can heat up every layer up to 180°C.

Metal powder: Bronze powder has been the choice of metal powder in this research because of its relative ease of sintering and because of the popularity of bronze parts. Sintered parts with fully alloyed bronze powders showed the best mechanical properties and the least shrinkage. Spherical shaped fully alloyed bronze powder with 44 micron particle size was used in this study.

Inhibitor selection: In this study, metal based salts have been chosen as the best inhibitor candidates, since salts are usually inexpensive and safe for both humans and the environment. The following important factors need to be considered in choosing an appropriate salt candidate: (1) the decomposition temperature; (2) molecular mass of the salt (3) Solubility. Accordingly, the best inhibitor was identified to be aluminum sulfate. The next successful candidate for inhibitor is sugar (C12H22011 ).

Sintering path: The sintering happens in an inert environment with Argon gas. Various sintering paths have been tested. A sintering path that yields reasonably small shrinkage is shown in Figure 4.

SIS Process with Aluminum Sulfate on Bronze Metal Powder:

Two sample parts manufactured by the SIS process are shown in figures 5 & 6. The layer thickness for the parts is .4 mm.

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Figure 4. Sintering path used in the experiments

Figure 5. A test bronze part can be separated from other portions after sintering

Figure 6. A test bronze part with internal holes

The properties of the bronze parts shown in figure 5 have been determined and listed tested in Table I.

Table 1. Mechanical properties of a bronze part built by the SIS process Shrinkage ( Linear shrinkage) Hardness ( Brinell Hardness) Ultimate Tensile Strength Density

6%-8% 60 30,000 psi 6.8 g W

During the SIS-Metal process the salt solution is first printed in the selected areas of metal powder. The decomposition of aluminum sulfate during the sintering process is shown as follows:

Al2 (SOA )3 -> Al2Oy + 3SO, & 2SO, -* 2S02 +02 (1)

In order to observe the mechanism of inhibition from sintering, different sections of the test bronze part were studied under the scanning electron microscope (SEM). The SEM micrographs are shown in figure 9.

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Figure 7. Scanning Electron Microscope (SEM) micrograph (a): a sample that is unsintered and without salt; (b): a sample that is sintered and without salt; (c): salt crystals are embedded in a

sample that is unsintered and with salt; (d): metal powder particles are covered by ceramic particles

During the SIS-Metal process the salt solution is first printed in the selected areas of metal powder. The decomposition of aluminum sulfate during the sintering process is shown as follows:

Al2 (SO, )3-+Al1Oi+ ISO, & ISO, -> 2S02 +02 ( 1 ) As it can be seen from figure 8, the energy-dispersive X-ray spectroscopy (EDS) analysis also shows the füll decomposition of aluminum sulfate. The decomposed metal oxide coats the adjacent metal particles. It can be seen that the dispersion of second phase particles (Aluminum Oxide ) retards the sintering of metal powder (Bronze). As it has been shown in a similar study for sintering pure copper mixed with Aluminum Oxide (Johnson 1976) it can be shown similarly that the dispersed particles of a second phase metal pinned to the surface of metal particles exerts a retarding force on the surface of the metal particle as it tries to move during sintering and therefore it prevents the transport mechanism in the second stage of sintering.

Figure 8. The SEM image of a sintered sample and the related EDS analysis result

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SIS Process with Sucrose as Inhibitor:

In this experiment instead of aluminum sulfate sugar (water-soluble crystalline carbohydrates with a formula of C„H2„On) was used as an inhibitor.

Figure 9. A test bronze part built by the SIS process using sugar as inhibitor

The decomposition of sugar starts in the printed sections starts at 185 °C and fully decomposes at around 820°C which is the maximum sintering temperature. The decomposition reaction is shown below:

CnH11On+x01-*\2H10 + yC + zC02 (2)

During the sintering process when the temperature raises the sugar coating decomposes into carbon, carbon monoxide and water where the carbon left behind gives a carbon coating to the affected metal particles. As it can be seen from the SEM micrographs in figure 10, Carbon particles are spread uniformly on the metal powder and it slows down the surface diffusion during the sintering process.

Figure 10. The SEM pictures of bronze sample printed with sucrose before sintering (left) after sintering (right)

SIS Process with Titanium Powder:

After finding bronze metal powder successful in the SIS process the next step was to try titanium parts which are in high demand in aerospace industries. The powder that was selected is titanium powder with -100/325 mesh size. The first step was to find the optimum sintering temperature. The titanium parts were sintered at 1250°C for 45 minutes in an inert (Argon gas) environment with a heating rate of 15°C/min.

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Both sucrose and aluminum sulfate with the same composition that was used for bronze powder was used for titanium powder as an inhibitor. The SEM pictures of the titanium powder with and without salt/sucrose, and also before and after sintering are shown in the figures below. Based on these SEM pictures it can be observed that the decomposition of salt crystals and sucrose plays a big role in preventing metal particles from being properly fused in the course of sintering.

As it can be seen from the SEM micrographs in figure 10 and EDS analysis, the same remarks and conclusions that were made for bronze parts regarding the reasons behind inhibition of sintering are also valid for titanium powder.

Figure 15. SEM images for titanium metal powder

Conclusion

The SIS-Metal process based on microscopic mechanical inhibition has been investigated for two different metals and two different inhibitors. Our experimental results have demonstrated that the SIS-Metal process can fabricate metal parts with desirable properties. Based on the experimental results an inhibition mechanism for preventing metal particles from sintering has been identified and described. The inhibition mechanism is based on the decomposition of salt and subsequent transition to ceramic particles (metal oxides) which covers the metal particles throughout the sintering process. By establishing the basics of the process this paper guides future researchers toward process improvement through attainment of higher accuracy and fabrication speed, and investigation of applicability of the process to other metal powder materials such as stainless steel.

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References

1. Mahdi Yoozbashizadeh, Behrokh Khoshnevis and Nozar Mozaka, " Development of Selective Inhibition Sintering (SIS) for Metallic Parts Fabrication", Internationaljournal of Advanced Manufacturing Systems, 13 (1)(2011), 107-122.

2. B. Asiabanpour, K. Palmer and B. Khoshnevis, "An experimental study of surface quality and dimensional accuracy for selective inhibition of sintering", Rapid Prototyping Journal, 10 (3) (2003), 181-92.

3. B. Asiabanpour, K. Palmer and B. Khoshnevis "Advancements in the selective inhibition sintering process development", Virtual and Physical Prototyping Journal, 1 (1) (2006), 43-52.

4. U. Behrendt and M. Shellabear, "The EOS Rapid Prototyping Concept", Computers in Industry, 28 (1995).

5. RM. German, Powder metallurgy science (Metal Powder Industries Federation, Princeton, 1994), 145-195.

6. M. Greul, T. Pintat and M. Greulich, "Rapid prototyping of functional metallic Parts", Computers in Industry, 28 (1) (1995), 23-28.

7. S. Kalpakjian and S. Schmid, Manufacturing Engineering and Technology (Pearson Prentice Hall, 5th

edition, 2006).

8. B. Khoshnevis, B. Asiabanpour, M. Mojdeh and K. Palmer, "SIS - a new SFF method based on powder sintering", Rapid Prototyping Journal, 9 (1) (2003), 30-36.

9. Y. Pelovski et al., "The thermal decomposition of aluminum sulfate in different gas phase environments", Thermochimica Ada, 205 (1992), 219-224.

10. E. Sachs et al., "Three dimensional printing: a candidate for the production of powder metal parts," Advances in Powder Metallurgy and Paniculate Materials, 1 (2000), 139-152.

11 R.L. Smith and S.V. Yanina, "Inhibition of sintering and surface area loss in phosphorus-doped corundum derived from diaspore", Journal of Am. Ceram. Soc, 85 (9) (2002), 2325-2330.

12. K.H. Stern, High temperature properties and thermal decomposition of inorganic salts with oxyanions (CRC Press, Ltd, 2001), 74-75.

13. H. Tagawa, "Thermal decomposition temperatures of metal sulfates", Thermochemica Acta, 80 (2001), 23-33.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

NUMERICAL SIMULATION OF COLD PRESSING OF ARMSTRONG CP-TITANIUM POWDER

Adrian S. Sabau1, Sarma B. Gorti2, William H. Peter', Wei Chen1, and Yukinori Yamamoto' 'Materials Science and Technology Division

2Computer Science and Mathematics Division Oak Ridge National Laboratory

One Bethel Valley Road, Oak Ridge, TN 37831-6083

Abstract Numerical simulation results for the cold pressing of Armstrong CP-Ti powder are presented. The computational model was implemented in the commercial finite element program ABAQUS™. Several simulation cases were conducted for cylindrical samples with different friction coefficients and different compaction pressures, under both single-action and dual-action uniaxial pressing. Numerical simulation results for the density distribution are compared against experimental data in order to validate the computational model.

Notice: This manuscript has been authored by UT-Battelle, LLC, under Contract No. DE-AC05-00OR22725 with the U.S. Department of Energy. The United States Government retains and the publisher, by accepting the article for publication, acknowledges that the United States Government retains a non-exclusive, paid-up, irrevocable, world-wide license to publish or reproduce the published form of this manuscript, or allow others to do so, for United States Government purposes.

Introduction Due to their low density, high specific strength and stiffness, and adequate creep and fatigue strength [1-7], titanium and titanium alloys could be used in many industrial and commercial applications. Despite its advantageous properties, Ti is underutilized in the industry due to the high cost of Ti that is produced by conventional reduction and melt processing technologies [8-10]. Recently, new lower cost titanium and Ti alloy production methods, e.g., the Armstrong Process technology, have been developed [11]. These processes require less energy than conventional reduction technologies while providing commercially pure titanium and titanium alloy powders that are suitable for powder metallurgy (PM) consolidation. The development of PM technologies that use the new low-cost titanium powders is required for producing economical titanium components with significant energy savings in both the titanium production and downstream industrial applications.

The consolidation process was recently studied for the near-net-shape manufacturing of the Ti alloy powders made by the Armstrong process [12]. It was demonstrated that more than 96% of theoretical density could be achieved by conventional pressing and sintering of the powder. SEM secondary electron (SE) images of as-received CP-Ti powders that were shown in Chen et al. [13] revealed that the Armstrong powder consisted of large agglomerates of several hundred microns in size. These agglomerates are made of fine primary particles of sizes less than 10 urn. During pressing, the powder compacts made of agglomerates are expected to deform in a different manner than powder compacts made of spherical solid particles. In order to reduce and eliminate defects in components processed via powder metallurgy, the powder compaction process must be designed to account for these Ti agglomerates.

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This paper deals with numerical simulation of the cold compaction, which is the first step in a powder compaction cycle. The aim of this work is to validate a computational model for the consolidation of Armstrong CP-Ti via powder metallurgy. Numerical simulations provide a crucial tool for understanding the non-uniform distribution of the density within the powder compacts. Cold pressing of metallic powder has been studied extensively using computer simulation models in order to understand the density non-uniformity in the PM components at different stages in their production. In order to deal with the sponge-like morphology of the powder, a model was developed that would handle a variety of yield functions, allowing greater flexibility for using various constitutive equations [14]. Aravas [15] and Zhang [16] have presented solution algorithms for obtaining the tangent/linearization moduli, which are consistent with the return-mapping algorithm. These solution algorithms were formulated for yield functions that have a general form, offering great flexibility to the model. The solution algorithms presented by Aravas [15] and Zhang [16] have been used extensively in many studies [17-23] and were implemented in Sabau et al. [14], where numerical simulations were performed for the cold pressing process of Armstrong CP-Ti powder.

Experimental data for the density distribution within cylinders after cold compression are presented in the next section. The section following that contains a succinct presentation of models used for cold pressing. Most of the mechanical property data, such as Young's modulus and Poisson's ratio as a function of relative density, which are necessary for the simulation of the deformation of powder compacts during cold pressing, were obtained from Sabau et al. [24], In the fourth section, numerical simulation results are presented for the cold pressing of samples under different conditions. The numerical simulation results for the average density are in good agreement with the experimental data, showing the capability of the model in predicting the distribution of relative density during cold pressing.

Experimental procedures and measurements The CP grade titanium powder used for this study was supplied by Cristal US Inc. and International Titanium Powder, Inc (ITP). This powder is produced using the Armstrong Process. Presently this powder is produced in a small batch reactor. The as-received powder was ball milled in ethanol in a 4 liter HDPE bottle containing 1 cm diameter high-density zirconia media (Glenn Mills) for 90 min. The powder was then dried at room temperature and the powder was sieved to remove any agglomerates that were greater than 40 mesh in size. The sample processing consisted of two steps: cold pressing and sintering. For the first step, a Carver Model "C," 24 ton hydraulic press was used along with a custom steel (D2) die. Samples were cold pressed uniaxially at 172 MPa (25 ksi) or 517 MPa (75 ksi) under single-action mode or dual-action mode. In single-action mode, the punch on one side moves, while the punch on the other side is held stationary. In dual-action mode, both punches move towards each other to press the powder. The experimental data on sample weight, dimensions, and relative density are shown in Table I. A typical microstructure picture is shown in Figure 1 for a sample that was pressed at 690 MPa (100 ksi).

Elastoplastic Constitutive Model For Cold Pressing

The constitutive response of the material during cold pressing is modeled using a pressure dependent elastoplastic model under the assumption of small elastic strains, following the approach detailed by Aravas [15] and Zhang [16].

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Sample ID Dl D2

SI

S2

Table I. Sample dimensions and density measurements after cold pressing Pressure

[ksi]

25 7.5 25

75

Uniaxial Pressing

Dual-action Dual-action

Single-action Single-action

Weight (g) 1.99 2.02

2.01

2.02

Diameter (cm) 1.275 1.271 1.275

1.273

Final Height (cm) 2.813 2.075 3.104

2.294

Density (g/cra3)

2.51 3.41 2.27

3.09

RD [%1

55.68 75.84 50.41

68.56

Figure 1. Metallographic picture showing porosity for a specimen that was cold pressed at 100 ksi.

Thus, the total strain increment is decomposed additively into elastic and plastic strain increments as, de=de'-^ep. For the case of linear isotropic elasticity, the stress-strain relationship appears, as:

a-C< :e% and c ^ ^ + ^ - f c X ' ® 7 ) (1)

where C is the fourth-order elasticity tensor, K is the elastic bulk modulus, G is the elastic shear modulus, 3 is the fourth-order identity tensor and / is the second-order identity tensor. The yield function is given in terms of the first and second invariants, p and q, of the stress tensor, and a set of scalar internal state variables, such as relative density, H , ct = l, 2,...,n,

1 ■JF de" -dA-

, where 3 is the hydrostatic stress, and V z is the equivalent stress, and S~a+ pi j s ^ e deviatoric stress tensor. The plastic flow rule is given by

v da, where dA is a positive scalar and g is the flow potential. It is common practice to apply the associated flow rule to porous solids, which implies 8~<P. The elastoplastic model is completed by a set of rules to update the internal state variables, dH =h (de ,a,H ) The internal state variables were taken to be the microscopic effective plastic strain, sp, and the relative density, D, whose rate of change is given respectively by the equivalent plastic work and the conservation of mass:

Dcjjry D = -Dtr(èp) (2)

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where sp is die plastic deformation rate tensor. In order to describe the plasticity of CP-Ti, a yield condition developed by Govindarajan and Aravas [25] was used. Material hardening parameters were determined by trial and error to match the experimentally measured variation of density with compaction pressure [14]. A fully implicit backward Euler integration of the elastoplastic constitutive model outlined above was presented in Sabau et al. [14]. The model was implemented as a user material subroutine (UMAT) in die finite element program ABAQUS.

Numerical Simulation Results Assuming axi-symmetry, a 2-D mesh for half of the cylinder cross-section was generated to simulate uni-axial cold pressing. The sample was assumed to be in contact with rigid bodies on the outer diameter, and on die top and bottom surfaces. A friction coefficient was imposed on all contact surfaces between the die and powder compact.

Simulations were carried out for cold pressing of Ti samples under bodi single-action and dual-action pressing. For the single-action case, die top punch was subjected to an applied pressure of 172 MPa (25 ksi) or 517 MPa (75 ksi) while die bottom punch was fixed. For the dual-action case, both me top and bottom punches were subjected to die applied pressure._The die was made of AISI D2 tool steel. Table II lists the properties used for me Ti powder and me steel die.

Table II. Material properties used in the cold pressing simulations Property Elastic modulus of Ti Poisson's ratio of Ti Initial yield stress of Ti Hardening coefficient for Ti Initial relative density of Ti Elastic modulus of tool steel Poisson's ratio of tool steel

Value 111 GPa(16,100ksi) 0.327 289.6 MPa (42 ksi) 124.1 MPa (18 ksi) 0.33 200 GPa (29,007 ksi) 0.28

Simulations were carried out for two values of friction coefficient between sample and die, 0.2 and 0.3. The initial height of die sample was computed based on me initial relative density and die measured weight and diameter of sample, as h0 = weight/(area * initialdensity). Tables III and IV summarize die data from die experimental measurements and die simulations. The results show that increasing die friction coefficient leads to a lower value for die average relative density. The simulated values are in fairly good agreement with the experimental measurements.

Table III. Computed results for cold pressing of CP-Ti samples using a friction coefficient of 0.2

Sample ID Dl D2 SI S2

Exp. Final Height (cm)

2.813 2.075 3.104 2.294

Comp. Final Height [cml 2.807 2.154 2.924 2.233

Exp. RD I%1

55.68 75.84 50.41 68.56

Comp. RD [%]

56.51 72.96 54.22 70.49

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Table IV. Computed results for cold pressing of CP-Ti samples using friction a coefficient of 0.3

Sample ID Dl D2 SI S2

Exp. Final Height (£m) 2.813 2.075 3.104 2.294

Comp. Final Height

2.947 2.262 3.136 2.398

Exp. RD

55.68 75.84 50.41 68.56

Comp. RD |%]

54.14 70.09 50.81 66.27

Figures 2 through 5 show the distribution of the relative density plotted on the deformed mesh for the different cases simulated. As mentioned earlier, these figures show half the cross-section of the cylindrical samples, where die left edge corresponds to the central axis of the sample while the right edge corresponds to the outer diameter (OD). Under dual-action pressing (Figures 2-3), the relative density is highest at the top and bottom corners along the OD while the lowest value is at the mid-thickness along the OD. In the case of single-action pressing, the highest relative density is at the top corner while the lowest value is at the bottom corner along the OD. An increase in the coefficient of friction leads to a lower average relative density. One noteworthy aspect is that this change comes from a decrease in die minimum value for the relative density for a given cold pressing case. As seen from the contour plots, there is little change in the maximum value of the relative density when the friction coefficient is changed. This suggests Üiat the maximum value of the relative density is governed to a larger extent by die applied pressure, while the minimum value is more sensitive to the friction coefficient.

Figure 2. Contours of relative density for sample Dl, dual-action pressed at 25 ksi, obtained from numerical simulations using a friction coefficient of (a) 0.2 and (b) 0.3.

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Figure 3. Contours of relative density for sample D2, dual-action pressed at 75 ksi obtained from numerical simulations using a friction coefficient of (a) 0.2 and (b) 0.3.

Figure 4. Contours of relative density for sample SI, single-action pressed at 25 ksi, obtained from numerical simulations using a friction coefficient of (a) 0.2 and (b) 0.3.

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Figure 5. Contours of relative density for sample S2, single-action pressed at 75 ksi, obtained from numerical simulations using a friction coefficient of (a) 0.2 and (b) 0.3.

Conclusions Results from the numerical simulations of the cold pressing of Armstrong CP-Ti powders were obtained. The computational model was implemented in the commercial finite element program ABAQUS™. Several simulation cases were conducted for cylindrical samples with different friction coefficients and different compaction pressures, for both single-action and dual-action uniaxial pressing. Numerical simulation results for the density distribution are compared against experimental data in order to validate the computational model. The results show that the highest relative density is governed to a larger extent by the compaction pressure, while the lowest relative density is influenced by the friction coefficient. The numerical simulation results for the average relative density are in good agreement with measured data.

Acknowledgements This research was sponsored by the U.S. DOE, and carried out at ORNL, under Contract DE-AC05-00OR22725 with UT-Battelle, LLC. This research was sponsored by the U.S. DOE, EERE Industrial Technology Program Office, under CPS Agreement #17881.

References 1. R. Boyer, G. Welsch, and E.W. Collings, Materials Properties Handbook: Titanium Alloys, 1994,

ASM international, Materials Park, OH. 2. G. Lutjering, J.C. William, Titanium, 2007, Springer, 1-14. 3. J. S. Montgomery et al., JOM, 1997, vol. 49, no. 5, 45^t7. 4. B.E. Hurless et al., The AMI'TIAC Quarterly, 2002, vol. 6, no. 2, 3-9. 5. China's Impact on Metals Prices in Defense Aerospace, Department of Defense, Office of the Deputy

Under Secretary of Defense (Industrial Policy), December 2005. 6. S. Luckowski, "The Application of Titanium in Army Armament Systems," Presentation at TMS

2003 Meeting, 2003. 7. S. Froes, Kirk-Othmer, Encyclopedia of Chemical Technology. Titanium and Titanium Alloys, 2001,

John Wiley & Sons, Inc.

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8. N. Savage, "A Cleaner, Cheaper Route to Titanium," Technology Review: An MIT Enterprise. 2006. http: www.lechnolosvreview.com printer friendly article.aspx?id 16963

9. A.D. Hartman et al. JOM, 1998, vol. 50, no. 9,16-19. 10. H. Friedrich et al., "Titanium in Automotive Applications -Nightmare, Vision, or Reality,"

Proceedings Iff' World Conference on Titanium, edited by G. Lutjering, Wiley-VCH, Weinheim, Germany, 2003, 3393-3402.

11. G. Crowley, AM&P, 2003, vol. 161, no. 11, 25-27. 12. Y. Yamamoto, JO. Kiggans, M.B. Clark, S.D. Nunn, AS. Sabau, and W.H. Peter, "Consolidation

Process in Near Net Shape Manufacturing of Armstrong CP-Ti/Ti-6Al-4V Powders," Key Engineering Materials, 2010, vol. 436, 103-111.

13. W. Chen, Y. Yamamoto, and W.H. Peter, "Investigation of pressing and sintering processes of CP-Ti powder made by Armstrong Process", Key Eng. Mat., 2010, vol. 436, 123-130.

14. A. S. Sabau, Sarma B. Gorti, William H. Peter, and Yukinori Yamamoto, "Process Simulation of Cold Pressing of Titanium CP-Ti Powders Produced by the Armstrong Process," in review for submission to the Powder Technology journal, 2011.

15. N. Aravas, "On the Numerical Integration of a Class of Pressure Dependent Plasticity Models,"/«/. J. Numer. Meth Engng., 1987, vol. 24, 1395-1416.

16. Z. L. Zhang, "Explicit Consistent Tangent Moduli with a Return Mapping Algorithm for Pressure-dependent Elastoplasticity Models," Comput. Methods Appl. Meek Engr., 1995, vol. 121,29—41.

17. M. A. Keavey, "A simplified canonical form algorithm with application to porous metal plasticity," Inter. J. Num. Meth. Eng., 2006, vol. 65, 679-700.

18. M. A. Keavey, "A canonical form return mapping algorithm for rate independent plasticity," Inter. J. Num. Meth. Eng., 2002, vol. 53, 1491-1510.

19. K. T. Kim and H. Park, "Effect of ceramic ball inclusion on densification of metal powder compact," Materials Science and Engineering A, 2000, vol. 282, no. 1-2, 29-37.

20. K. T. Kim and H. C. Yang, "Densification behavior of titanium alloy powder during hot pressing." Materials Science and Engineering A, 2001, vol. 313, no. 1-2,46-52.

21. Y. S. Kwon, H. T. Lee, etal, "Analysis for Cold Die Compaction of Stainless-Steel Powder,' 'Journal of Engineering Materials and Technology, 1997, vol. 119, no. 4,366-373.

22. S. C. Lee and K. T. Kim (2002). "Densification behavior of aluminum alloy powder under cold compaction," International journal of Mechanical Sciences, 2002, vol. 44, no. 7, 1295-1308.

23. S. C. Lee and K. T. Kim, "A study on the Cap model for metal and ceramic powder under cold compaction," Materials Science and Engineering A, 2007, vol. 445-446, 163-169.

24. A. S. Sabau, J. O. Kiggans, W. H. Peter, and D. L. Erdman III, "Material Properties for the Simulation of Cold Pressing of Titanium Armstrong CP-Ti/Ti-6Al-4V Powders," 2010 International Conference on Powder Metallurgy & Paniculate Materials, PowderMet 2010, June 27-30, Hollywood, FL.

25. R. M. Govindarajan and N. Aravas, "Deformation Processing of Metal Powders: Part I—Cold Isostatic Pressing," Int. J. Mech. Sei., 1994, vol. 36, no. 4, 343-357.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

THE EFFECT OF COKE PARTICLE SIZE ON THE THERMAL PROFILE OF SINTERING PROCESS PRODUCT

N. Tahanpesaranedezfuly, A. Heidary Moghadam Department of Metallurgy, Azad University,Dezfool Branch

P.O.Box 133, University Blvd. Dezfool-Iran

Keywords: sintering process, coke particles size, permeability, thermal distribution

Abstract In this research, the effect of coke particle size on thermal profile in the production process is

investigated in the laboratory environment. Baking process and sinter production was performed for different ranges of coke particle size while other parameter like, iron ore, lime and sintering mixture particle sizes, suction conditions and so on were kept constant. Thermal profile was achieved by measuring the temperature in different places of sintering bed. Investigating the thermal profile in different tests confirmed that coke grain size has a significant effect on thermal profile. The effect of coke particle size on coke combustion velocity, maximum temperature, the amount of produced melting phase as well as the flame front speed was understood by analyzing the achieved profiles and consequently the optimal grain size was known. The coke suitable particles size could have a great effect on quality of the produced sinter.

Introduction Iron ore sintering process, which combustion of coke fines happens in the materials mixture

bed (consisting of iron ore, lime, coke and so on) by air suction (from the bed surface in to it) is known as an important process for converting the iron ore to iron. This technique was innovated in the late nineteenth century. By 1911 and as a consequence of equipping the sintering machines to conveyer belt, in addition to tremendous increase in efficiency of the sinter production in these machines, considerable development was made in improving the blast furnace process. It is well confirmed by experience that a very change in the controlling parameters of chemical and physical properties of the sintering process product, could greatly affect the blast furnace production efficiency in different aspects, qualitative and quantitatively especially its energy consumption.

Generally the physical and chemical properties of the sinter could be concise to three different aspects: mineralogical, structurally and its porosity. The most effective parameters on controlling these three topics or on produced properties of the sinter are: granules size and composition, composition and relative properties of mixture components (iron ore, flux, coke), mineralogical type and structure of mixture components, thermal profile of the process, amount of water of the mixture, density of the mixture before baking process and size distribution of mixture components [1-5].

Among the parameters, the thermal profile of the process along the bed has significant effect on quality of the produced sinter, because the cooling rate of the baked sinter has a considerable effect on crystallizing the binder phases and thermal stresses as well. The latter greatly affects the sinter properties. In addition, produced thermal profile along the bed is greatly dependent on the behavior of the combustion zone. Furthermore, the behavior of combustion zone is intensively affected by coke distribution and its particles size. Coke is the major energy producer in this process. Properties like porosity and diffusivity, strength, reducibility of the sinter are affected by coke particles size and distribution in sintering mixture. Existence of big coke

529

particles in sintering mixture leads to a heterogeneous charge and thermal profile. Furthermore, temperature near the big particles, increases and consequently it might cause charge melting and decreasing the product diffusivity. On the other hand, small coke particles burn faster and as a consequence the thickness of the layer with highest temperature decreases. It's worth saying that very small coke particles, decrease the diffusivity of the bed.

In recent years, the effect of coke particles size has been interesting to many investigators, which some of these studies are shown in table I [6-7]. The difference in the results of these researchers is because of difference in their tests conditions including size, the type of consuming material and so on.

Although there are considerable researches on the sintering process, however, these works are focused on the sinter quality. There is little emphasis on the effect of coke particle size on the thermal profile of the sintering bed. In this research, which raw materials were from Isfahan Iron Melting Company, the effect of coke particles size on thermal profile in sinter production process was investigated in laboratory and optimal coke particle size in the same baking conditions was specified.

Table I Optimum coke particle size obtained in previous researches

Researcher

Coke particle size

Kadama et al. (1982)

0.25-3 mm

Peter et al. (1985)

1-3.5 mm

Barbosa et al. (1985)

0.3-3.36 mm

Loo et al. (1991)

1.5-2.5 mm

Experimental method Eight different coke particles sizes were chosen and were studied. The first particle size was

considered as the reference sample or standard sample. The other particle sizes were chosen to study the effect of coke big particle size as well as its small particles. Chosen particle size and percentage of different ranges of coke particles are shown in table II. Each sample was denoted by a capital letter. Sintering was performed for each range of these in laboratory after specifying the coke particle size in different tests.

For baking process, first a mixture of iron ore, return sinter, burnt lime, raw lime and coke were mixed in dry condition for five minutes, then the granulating process was performed. The produced mixture was baked in a sinter production pot. Temperature changes along the bed were measured by three thermocouples and were recorded by computer software as well. Figure 1, illustrates the sinter production machine and where thermocouple locations. Furthermore, the results of the tests were proved by doing the experiments in five sequences.

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Table II Percentage of different ranges of coke particles

Sample nun

-0.09 0.09-0.125 0.125-.212 0.212-0.25 0.25-0.335 0.335-0.5 0.5-0.71 0.71-1 1-1.4

1.4-1.7 1.7-2.88

2.88-3.35 3.35-6.3 6.3-9.65 +9.65

A Standard sample 8.49

4 4

6.7 7.58 5.59 6.5

7.55 7.32 7.31 7.44 8.39 15

3.32 0.22

B 0-3.35

10.48 4.95 4.94 8.27 9.35 6.9 8.02 9.31 9.04 9.02 9.18 10.36

---

C 0-2.88

11.73 5.13 5.53 9.25 10.46 7.71 8.97 10.41 10.11 10.09 10.27

----

D 0.212-2.88

---

9.25 10.46 7.71 8.97 10.41 10.11 10.09 10.27

----

E 0.212-3.35

---

10.42 11.79 8.69 10.11 11.73 11.39 11.37 11.57 13.05

---

F 0.355-3.35

-----

11.12 12.94 15.02 14.58 14.55 14.81 16.71

---

G 0.5-3.35

------

14.62 16.97 16.47 16.44 16.73 18.88

---

H 0.71-3.35

-------

19.83 19.25 19.21 19.55 22.06

---

Figure 1. Schematic of sintering pot and thermocouples

Result and discussion Different parameters including coke combustion rate the achieved maximum temperature, the

amount of molten phase formation and flame front speed were extracted by analyzing the thermal profiles of different tests, then the effect of coke particle size was specified. Figure 2, shows the sinter thermal profiles of the all samples. Figure 3, schematically shows the studied parameters.

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Figure 2. Thermal profiles for samples A-H

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Figure 3. Schematic of thermal profile and studied parameters a) Flame front speed: velocity of flame front between two thermocouples,

b) T max: maximum temperature, dT/dt: coke combustion rate, liquidus is specified asll64*C

Coke combustion rate The meaning of coke combustion rate in its definition is temperature increasing slope, in

which this slope is shown in figure 3. Because the coke is the only parameter which produce energy in sinter process, so the slope of these diagrams depends on coke reactivity and existent oxygen amount for burning the coke.

By decreasing the size of coke particles, its specific surface increases and so its reactivity and burning ability as well. Figure 4, shows changes in the coke burning rate for the test samples. As it's illustrated, the maximum combustion rate is accessible by expelling the particles bigger and smaller than 0.212mm and 3.35mm, respectively. As it was mentioned earlier, coke combustion rate depends on coke reactivity and the amount existent oxygen. Providing the necessary oxygen depends on diffusivity of the charge bed. So diffusivity of charge bed and coke reactivity is two important and effective parameters governing the coke combustion rate. Coke reactivity

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increases by eliminating the coke big particles, but diffusivity of the charge bed decreases. Otherwise, diffusivity of the bed improves by expelling the small coke particles and coke reactivity decreases, these optimal conditions are well shown by sample E. however, coke reactivity decreases by continuing the elimination of coke small particles. Loo et al. [1-3] showed that diffusivity of the bed could be improved by choosing a suitable range.

Figure 4. The effect of coke particles size on coke combustion rate

Maximum temperature The amount of the produced heat per time and the maximum temperature is controlled by

diffusivity and reactivity of the coke, which in turn are affected by the coke particle size. The amount of heat produced per time increases by changing the coke particle size in sinter process and choosing a suitable range of coke particle sizes as well, because the diffusivity and coke reactivity increases. It is because in one layer, the coke particles simultaneously and rapidly become active and as consequently, heat produces concentrate in flame front and finally the temperature of flame front increases. Figure 5, shows a comparison of maximum temperature of the second thermocouple in the samples. As it can be seen, the sample E has the maximum increase in temperature among the samples. Enhancement of the coke reactivity and bed diffusivity parameters by choosing this range of coke particle size is the reason for that.

Figure 5. The effect of coke particle size on maximum temperature

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The amount of molten phase According to raw materials composition used in these tests, starting temperature of molten

phase formation was specified as 1164°C [8]. Knowing the maximum temperature and its related composition, the amount of molten phase formation was specified by using the Lever rule. These results are shown in figure 6. as it can seen, the maximum amount of molten phase formation is for sample E. the amount of molten phase formation directly depends on the maximum temperature of flame front and as it mentioned earlier, it has a great dependence on coke reactivity and diffusivity of charge bed. Kasai et al. [4-5] showed that as sintering temperature increases, the amount of molten phase and its fluidity increases.

Figure 6.The effect of coke particle size on liquid fraction

Flame front speed Velocity of flame front between the first and third thermocouples in the test samples are

compared in figure 7. As it is clear by eliminating the coke big particles, which have little activation and by eliminating coke small particles, which reduce the bed diffusivity, and choosing the right range, which has the optimal bed diffusivity and coke reactivity, velocity of flame front could be increased. This increase in the velocity could results to an increase in the maximum temperature of the flame front, because it increases the amount of produced heat per time. Sample E, which includes the coke particle range of 0.212 to 3.35 mm, shows the maximum velocity of the flame front and consequently the time for the full baking in this sample is the minimum among the samples.

Figure 7. The effect of coke particle size on flame front speed

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Conclusion In this research, the change in the coke particle size and its effect on temperature distribution along the bed of the sinter vessel (thermal profile) were investigated. Results showed that the range of 0.212-3.35 mm, is the optimal range, which has these consequences:

i) Increasing the coke combustion velocity and as a result, increasing the amount of produced heat per time. ii) Increasing the maximum temperature. iii) Increasing the molten phase formation. iv) Increasing the flame front rate along the bed.

In fact, eliminating particles smaller than 0.212mm and larger than 3.35mm, creates an optimal condition for bed diffusivity and coke reactivity and caused mentioned results. It's clear that, these results have a significant effect on the quality of the produced sinter including its strength, diffusivity and structure.

References

1. C. E. Loo, "A Perspective of Goethitic Ore Sintering Fundamentals", ISU International, 45 (2005) 436-448. 2. C. E. Loo, D. J. Wong, "Fundamental Factors Determining Laboratory Sintering Result", ISU International, 45 (2005), 449-458. 3. C. S. Teo, R. A. Mikka, C. E. Loo. "Positioning Coke Particles in Iron Ore Sintering" , ISU International, 32 (1992) 1047-1057. 4. E. Kasai, W. J. Rankin, J. F. Gannon, "The Effect of Raw Mixture Properties on Bed Permeability during Sintering", ISU International, 29 (1989), 33-42. 5. M. V. Ramos, E. Kasai, J. Kano, T. Nakamura, " Numerical Simulation Model of the Iron Ore Sintering Process Directly the Agglomeration Phenomenon of Granules in the Packed Bed " , ISU International,^ (2000), 448-454. 6. P. R. Dawson, "Research Studies on Sintering and Sinter Quality", Ironmaking and Steelmaking, 20 (1993), 137-143. 7. C. E.Loo, "Role of Coke Size in Sintering of a Hematite Ore Blend", Ironmaking and Steelmaking, 18 (1991), 33-40. 8-K. Higuchi, M. Naito, M. Nakano and Y. Takamoto, " Optimization of chemical composition and microstructure of iron ore sinter for low temperature drip of molten iron with high permeability" ISU, 44 (12) (2004), 2057-2066.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

POWDER MATERIAL PRINCIPLES APPLIED TO ADDITIVE MANUFACTURING

David L. Bourell and Joseph J. Beaman

Laboratory for Freeform Fabrication, Texas Materials Institute, Mechanical Engineering Department, University of Texas at Austin, 1 University Station, Austin, TX 78712, USA

Densification, Sintering, Laser, Sintering, Powder

Abstract

Laser Sintering (LS) is an additive manufacturing process in which a part is constructed from powders without the use of part-specific tooling. Production of metallic, ceramic and composite parts in some cases requires some form of pre-processing or post-process sintering to achieve full density. Powder Densification Maps are a tool for optimizing the pre- and post-processing parameters. Such maps are computational representations of part density as affected by time, temperature, pressure and materials properties. This paper summarizes LS developments with emphasis on the utility of powder densification mapping of powder pre-processing and part post-processing. Specific emphasis includes developments in Powder Densification Map production for zirconia, Ti-6A1-4V, nickel-based Alloy 625 and copper. A comparison of theoretically predicted and experimentally determined densities for a variety of processing conditions is presented.

Introduction

Laser Sintering

Laser Sintering (LS) is a manufacturing process in which a part is produced without the need for part-specific tooling [1-3]. It competes effectively with other manufacturing processes when part geometry is complex and the production run is not large. Traditionally, this was limited to prototype production, although tooling applications are now appearing.

A schematic of this process is shown in Figure 1. The process begins by first depositing a thin layer of powder into a container. The powder surface is raster-scanned with a laser beam. Beam intensity is modulated to rase the powder in areas to be occupied by the part at that particular cross section. In areas not sintered, the powder remains loose and may be removed once the part is completed. Successive layers of powder are then deposited and sintered until the entire part is produced. Each layer is sintered deeply enough to fuse it to the underlying layer.

This paper describes several applications of powder densification maps to advance solutions in direct LS of metallic and ceramic powders. The findings illustrated here have been published in a previous paper by the authors [4]. The LS process itself entails melting and resolidification of at least one component. Time-dependent plasticity issues arise in pre-processing of powder to make it suitable for LS

Figure 1. Schematic of the Laser Sintering Process.

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and in post-processing to obtain desired density.

Three applications are chosen. First, nanocrystalline, agglomerated ceramic powder is pre-sintered to density particulate within agglomerates without sintering the agglomerates. The second application involves post-processing by hot isostatic pressing (HIP) of LS metallic parts. The third application is post-process particle neck formation to produce a high-porosity metallic skeleton that can be infiltrated in a later step.

Time-Dependent Plasticity Issues in Laser Sintering

In all cases, solutions were sought by generation of Ashby Temperature Densification Maps [5-7]. These are computational plots of the relative density of a part as a function of temperature. Their generation is described in the next section.

Powder densification rate is dependent on a variety of factors including: intrinsic material properties such as yield strength and diffusion rates; the dominant densification mechanism; the initial density, part geometry, powder particle size and size distribution; and processing parameters such as temperature and pressure. The ability to model the densification process is beneficial and permits planning of more advanced process optimization such as control and modification of microstructure or geometric distortion.

The influence and interaction of most of these varying mechanisms is most readily represented in a graphical format across the entire processing envelope. This graphical representation has become known as a densification map. The basis of this concept was developed by Weertman [8]; however, the construction of the maps was brought to fruition by Ashby [5] in the early 1970's. The constitutive equations for the generation of the maps have been incorporated into a computer program which was used to create the deformation maps provided here.

Densification Modeling

Powder Densification

The models developed for both conventional pressureless sintering and HIP assume an initial structure consisting of uniformly sized, perfectly spherical particles. This assumption greatly simplifies the calculations in comparison to a system of arbitrary, irregularly shaped particles, and in fact, is not too far removed from the powders selected here for LS. Initially the spherical particles are assumed to be in point contact. As the sintering process proceeds, the contact area between particles increases to form "necks" or "bridges" between adjacent particles.

In early stages of sintering, particle necks provide strength, but very little densification takes place. As sintering proceeds, necks continue to enlarge, and particles begin to coalesce, resulting in increasing bulk density. The enlarged necks between particles produce an interconnected pore structure throughout the powder mass typified in Figure 2a. While this interconnected pore structure exists, it serves as an important conduit for mass transport; however, when a theoretical density of approximately

Figure 2. Illustration of (a) interconnected Stage I pore struc-ture and (b) isolated, Stage II pore structure From Ref. 9.

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92% is reached, the continuous structure becomes unstable and collapses to form isolated pores, as illustrated by black dots in Figure 2b. These pores during sintering become spherical in shape. This transition is defined as the boundary between the initial sintering conditions, Stage I, and the later Stage II process. This change in sintering behavior requires a modification of the equations used to described the densification process since the pore shape and connectivity affect densification rates. As a result, two different equations exist for each densification mechanism, one for Stage I and one for Stage II.

Driving Force for Densification and Densification Mechanisms

For pressureless sintering, a reduction in total surface area is the only driving force for the densification of the powder mass. In the case of HIP, an additional driving force stemming from the stress created by the compaction pressure is also active. To provide a convenient and unified form of the two driving forces, the surface area component of the driving force is converted to an effective stress and summed with the net applied HIP gas pressure, as shown in Equation 1 and Equation 2. A complete derivation of the equations is given by Ashby [10].

P . o . ^ P e x . - P o + S A 2 ^ ^ ^ ] [Stage I] (1)

1

(1-A )A y ( 6A ^3 P 1 0 „ = P „ , - r , n l ^ + 2 ^ _ ) [Stage„] (2)

where Ptotai = Effective Applied Driving Pressure, Pext = HIP Pressure, Po = Initial Pore Pressure, Pun = Internal Pore Pressure at Pore Close-Off ( Stage I to Stage II Transition), A = Relative Density = p/pth, p = Part Density, pth = Theoretical Density, y = Powder Surface Energy, R = Powder Radius, A0 = Initial Relative Density and Ac = Relative Density at Pore Closure/Isolation.

In practice, the contribution due to the surface energy reduction term is important in conventional, pressureless sintering, but it is minor relative to realistic hot isostatic pressing (HIP) pressures. For example, using fairly typical values of R = 25 um and y = 2 J/m2 it is seen that even near the end of the densification process (A = 0.97), where the surface energy contribution is greatest, the effective pressure is still only 0.93 MPa. With typical HIP pressures on the order of 100 MPa, it is apparent that there is a substantial driving force increase associated with HIP compaction relative to conventional pressureless sintering.

In all stages of sintering, several densification mechanisms act simultaneously. In Ashby's initial analysis for pressure sintering maps, four independent mechanisms were included: lattice diffusion, grain boundary diffusion, power-law creep (or dislocation creep), and plasticity (yielding/dislocation glide). In later work, the additional mechanisms of Nabarro-Herring creep and Coble creep were included for cases where the internal grain size of the particles is substantially smaller than the particle diameter [7]. Grain growth and grain boundary separation from pores during Stage II densification are also incorporated into the model. A complete listing of the data used in constructing the densification maps is given in the Table.

Experimental Results

Zirconia Pre-Processine

The first utilization of time-dependent plasticity in LS is pre-processing of zirconia. Zirconia is important for generation of molds for titanium metal casting. Zirconia is a material of choice, primarily due to its protection of the contacting titanium from formation of a surface,

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Table I - Material Property Values for Densification Map Calculations [From Ref. 10]

oxygen-rich alpha case [11]. Fine particulate zirconia is preferred from considerations of surface finish. Nanocrystalline (D = 24 nm) zirconia was available in polymer-bound, agglomerated form (50 um). This was unacceptable for LS processing because the binder would melt and flow in an unpredictable fashion during LS, and the coarse particles would not hold together. The need then was to determine a thermal excursion by which the nanocrystalline particulate might be sintered together to high density and by which the 50 um agglomerates would not sinter.

Pressureless temperature densification maps were generated for zirconia with a particle size of 24 nm and 50 (im. These appear as Figures 3 and 4. It is observed that zirconia fine particles sinter lightly in 2 hours at 1050° C and almost completely density in 2 hours at 1250°C. At 1250°C, the dominant densification mechanism is boundary diffusion initially which transitions to volume (bulk) diffusion at about 45% relative density. Further, from Figure 4, it is seen that 50 urn zirconia is quite resistant to pressureless densification to at least 2000"C.

Figures 5a and 5b show low and high magnification SEM micrographs of yttria-stabilized zirconia powder held 2 hours in air at 1050°C. In Figure 5b, the fine particulate is visible and has lightly sintered. Figures 5c and 5d are SEM micrographs of powder held 2 hours at 1250"C.

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Figure 3. Densification Map for 24 nm zirconia, reflective of the densification kinetics of zirconia agglomerate interiors. Symbols are experimental results for a 2 hour hold.

Figure 4. Densification Map for 50 (am zirconia, reflective of the densification kinetics of zirconia particles. The large particles are very resistant to pressureless densification.

The coarse particles are still unsintered, but the fine-particle interiors are completely densified. This is consistent with the densification maps.

Hot Isostatic Pressing (fflP^ of Alloy 625 and TJ-6A1-4V

The densification predictions of the HIP model were verified through a series of HIP experiments. Alloy 625 and Ti-6A1-4V powders were processed over a range of temperatures, pressures, and times. The density achieved for a particular set of processing conditions was measured by gas pycnometry or Archimedes' method and compared to the value predicted by the HIP maps. Powder was placed in borosilicate glass, evacuated to at least 5xl0"5 Torr, degassed at 350°C for a minimum of 24 hours to remove the majority of surface adsorbed gases as well as some species dissolved within the powder [12], and the glass tube was sealed prior to HIP. The glass capsules were processed in an ABB QIH-3 HIP system equipped with a graphite heating element. The process cycle began with a 10°C/min ramp from room temperature to the desired processing temperature. Once the operating temperature was reached, the pressure within the HIP vessel was

Figure 5. SEM Micrographs of Agglomerated Zirconia. (a) Held in air at 1050°C for 2 hr. (b) Held in air at 1050"C for 2 hr. Fine, low-density nanocrystalline structure present within 50 um particles, (c) Held in air at 1250°C for 2 hr. (d) Held in air at 1250°C for 2 hr. Within 50 urn particles, the nanocrystalline particulate has fully densified.

541

raised. After the desired processing time had elapsed, the temperature was decreased at 15°C/min and the pressure was gradually reduced to atmospheric pressure.

A HIP map was generated for nickel-based Alloy 625 as shown in Figure 6. All three operating parameters (time, temperature, and pressure) were varied in the HIP experiments. In most cases, the predicted values are within one or two percent of the experimentally determined density. For the time/temperature range of interest, power-law creep dominates initial densification. A transition to boundary/volume diffusion accommodation occurs at higher densities. Nabarro-Herring creep is dominant at the onset of Stage II densification.

Titanium alloy Ti-6A1-4V was LS processed to produce a gas-impermeable skin with internal loose and lightly sintered powder. The goal was to create a sample which could be containerless HIPped [13,14]. Here the question deals with optimization of the HIP cycle to produce parts with full density with optimum temperature/time parameters. The results are shown in Figure 7 for an applied HIP pressure of 69 MPa (10 ksi). The power-law creep data for titanium was taken from Warren, et al. [15]. The creep stress exponent was taken to be 3 with a creep activation energy of 200 kJ/mol. The agreement is good at lower temperatures, but the material appears to sinter faster than predicted at temperatures greater than about 750°C. It has been clearly demonstrated by a number of researchers that fine grained (<10 urn) Ti-6A1-4V exhibits superplastic behavior at temperatures between 750°C and 950C [16 -21]. These authors have suggested that a grain boundary sliding (GBS) mechanism accounts for the behavior of the material under the specified processing conditions. This mechanism is currently not included in the HIP map model. A number of constitutive relationships have been developed to describe deformation associated with grain boundary sliding (or shearing). The incorporation of an appropriate rate equation into the model could improve the prediction accuracy for superplastic materials such as Ti-6A1-4V.

Figure 6. HIP map for Alloy 625. HIP pressure Figure 7. HIP Map for Ti-6A1-4V. HIP = 69 MPa, particle size = 50 urn. pressure = 69 MPa, particle size = 55 um. Experimentally determined values have been Experimentally determined values have been plotted for 1 hour (•) and 8 hours (■). plotted for 0.1 hour (•), 2 hr (♦), and 4 hr (■).

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Neck Formation in Copper Powder

The third topic is neck formation in copper powder. The goal here is to make a LS copper preform with low relative density suitable for liquid metal infiltration. Copper has high reflectivity to visible and infrared wavelengths, so LS was accomplished indirectly. A polymeric powder was mixed with the copper, and this polymer melted and "glued" the copper particulate together during laser scanning in LS. A post-process burn-out took place to remove the polymer and to lightly sinter the copper via neck formation. The issue here is identification of the pressureless time-temperature profile necessary to burn out the polymer and to cause neck formation in copper without allowing too much copper densification.

Spherical copper powder of average particle size 45 um was chosen. A pressureless temperature densification map is shown as Figure 8. Densification initiates in about 5 hr at 800"C. Further, the dominant initial densification mechanism is shown to be boundary diffusion. In furnace trials, the copper particles formed necks but did not substantially density when given this thermal exposure, Figure 9, in agreement with the densification map.

Figure 9. SEM Micrograph of copper powder after sintering for 5 hr at 800°C. Neck formation is extensive

Figure 8. Densification Map for Copper. The powder but little overall densification is is resistant to pressureless densification. The data observed, consistent with the point is for a 5 hour hold at 800'C. densification map, Figure 8.

Summary and Conclusions

Understanding and application of time-dependent plasticity is crucial to advancing our knowledge in the area of Laser sintering. Ashby densification maps are an excellent tool for incorporating the effects of stress, time and temperature on both mass transport mechanisms and the resulting densification rate. Specific examples are powder pre-processing of zirconia, hot isostatic pressing of nickel-based Alloy 625 and Ti-6A1-4V, and part post-processing of copper as a preliminary step to part infiltration. Powder densification maps were constructed by calculating the densification contributions of several independent deformation mechanisms: lattice diffusion, boundary diffusion, power-law creep, Nabarro-Herring creep, and Coble creep. The density predictions were verified experimentally.

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References

1. D.L. Bourell et al., "Solid Freeform Fabrication: An Advanced Manufacturing Approach" (Paper presented at the 1st Intl Solid Freeform Fabrication Symposium, Austin, Texas, 6 August 1990), 1-7.

2. C.R. Deckard, "Method and Apparatus for Producing Parts by Selective Sintering", US Patent #4,863,538.

3. H.L. Marcus et al., "From Computer to Component in 15 Minutes: The Integrated Manufacture of Three Dimensional Objects", JOM, 42 (4) (1990), 8-10.

4. D.L. Bourell et al., "Powder Densification Maps in Selective Laser Sintering", Advanced Engineering Materials, 4 (9) (2002), 663-9.

5. M.F. Ashby, "A First Report on Sintering Maps", Ada Metallurgica, 22 (1974), 275-289. 6. F.B. Swinkels and M.F. Ashby, "A Second Report on Sintering Maps", Ada Metallurgica,

29 (1981), 259-281. 7. A.S. Helle, K.E. Easterling, and M.F. Ashby, "Hot-Isostatic Pressing Diagrams: New

Developments", Ada Metallurgica, 33 (12) (1985), 2163-74. 8. J. Weertman, "Discussion of An Empirical relation Defining the stress Dependence of

Minimum Creep Rate in Metals", Trans Metallurgical Society AIME, 227 (1963),1475-6. 9. Randall M. German, Powder Metallurgy Science, 2nd Edition. (Princeton, NJ: Metal

Powder Industries Federation, 1994), 245-255. 10. M.F. Ashby, Sintering and hostatic Pressing Diagrams, (Published by Author,

Department of Engineering Cambridge, England, 1990.) 11. E.D. Calvert, "U.S. Bureau of Mines Report of Investigations RI 8541", (Government

Printing Office, Washington, DC. 1981). 12. B. Engel and D.L. Bourell, "Titanium Alloy Powder Preparation for Selective Laser

Sintering", Rapid Prototyping Journal, 6 (2) (2000), 97-106. 13. S. Das et al., "Processing of Titanium Net Shapes by SLS/HIP", (Paper presented at the

9* Intl Solid Freeform Fabrication Symposium, Austin, Texas, 10 August 1998), 469-77. 14. Suman Das et al., "Producing Metal Parts with Selective Laser Sintering/Hot Isostatic

Pressing", JOM, 50 (12) (1998), 17-20. 15. J. Warren, L.M. Hsiung, and H.N.G. Wadley, "High Temperature Deformation Behavior

of Physical Vapor Deposited Ti-6A1-4V", Ada Metall. Mater., 43 (7) (1995) 2773-87. 16. D. Lee and W.A. Backofen, "Superplasticity in Some Titanium and Zirconium Alloys",

Trans Metallurgical Society AIME, 239 (1967),1034-40. 17. A. Arieli and A. Rosen, "Superplastic Deformation of Ti-6A1-4V Alloy", Metall Trans,

8A (1977), 1591-96. 18. A.K. Ghosh and C.H. Hamilton, "Mechanical Behavior and Hardening Characteristics of

a Superplastic Ti-6A1-4V Alloy", Metall Trans, 10A (1979), 699-706. 19. N.E. Paton and C.H. Hamilton, "Microstructural Influences on Superplasticity in Ti-6A1-

4V", Metall Trans, 10A (1979), 241-50. 20. J.A. Wert and N.E. Paton, "Enhanced Superplasticity and Strength in Modified Ti-6A1-

4V Alloys", Metall Trans, 14A (1983), 2535-44. 21. M.L. Meier, D.R. Lesuer, and A.K. Mukherjee, "a Grain Size and ß Volume Fraction

Aspects of the Superplasticity of Ti-6A1-4V", Mali Sei and Engg, A136 (1991), 71-8.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

PROCESSING CHALLENGES OF DUAL-MATRIX CARBON NANOTUBE ALUMINUM COMPOSITES

A.M.K. Esawi1, K. Morsi2,1. Salama1, H. Saleeb1

'Department of Mechanical Engineering The Yousef Jameel Science and Technology Research Center (YJSTRC)

The American University in Cairo (AUC) Cairo, Egypt

department of Mechanical Engineering, San Diego State University

5500 Campanile Drive San Diego, CA 92182, USA

Keywords: Carbon Nanotube reinforced composites, Aluminum-CNT dual matrix composites

Abstract

The interest in nanostructured materials has grown considerably in recent years. Significant enhancements in strength have been reported. A common problem, however, is the associated reduction in the materials ductility. Efforts to overcome this ductility challenge by designing multi-modal or hierarchical microstructures have recently been reported. In this work, we report on the processing of dual matrix carbon nanotubes composites in which composite particles of aluminum reinforced with carbon nanotubes are embedded within a soft aluminum matrix. Such approach aims at combining the high strength of the CNT-reinforced region with the ductility of the soft aluminum matrix. Preliminary results, however, show that the interface quality is pivotal in promoting bonding between the two dissimilar particles since a poor interfacial bond is found to lead to deterioration in the composite properties; thus making it impossible to achieve the desired paradox of strength and ductility. Processing-related challenges in this regard are discussed.

Introduction

Due to their exceptional properties, carbon nanotubes (CNTs) have been recently used to reinforce materials in order to enhance their mechanical, electrical and thermal properties. The interest in CNT- metallic composites has been growing significantly over the past decade. In particular, carbon nanotube reinforced metal composites have been under intense investigations with the goal of generating composite materials with enhanced properties [1-3]. The consolidation of CNT-aluminum powder-based composites has so far involved processes such as hot pressing, hot extrusion, powder rolling, high pressure torsion, cold spraying, spark plasma sintering (SPS), and spark plasma extrusion (SPE) [4-13]. Most work has focused on the generation of composite materials with homogeneously dispersed carbon nanotubes. Nearly all the research groups reported enhanced mechanical properties to various degrees but this was generally combined with the undesirable loss in ductility.

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The dual matrix composite microstructural design, where the matrix is selectively reinforced in localized regions within the microstructure separated by ductile unreinforced matrix, may present significant benefits. This unique microstructural design has been documented to enhance both the wear resistance and toughness over conventional (single matrix) composites [14-16]. Other potential benefits include an ability to tailor the properties of the final composite which may give rise to controlled properties and enhanced formability.

A recent paper by two of the authors addressed the spark extrusion of single and dual matrix CNT-aluminum composites [17]. This paper presented for the first time CNT-aluminum composites based on the dual matrix design. However, the paper focused on the effect of the dual matrix design on the formability of the composite using spark plasma extrusion. No investigation of either the tensile behaviour of this novel material or its fracture toughness behavior was carried out. Furthermore, it has not yet been attempted to prepare the dual matrix composite using more conventional techniques such as compaction and hot extrusion; expected to provide more flexibility in preparing the unique sample geometry needed for fracture toughness tests.

Prior to adding the ductile matrix, the CNTs have to be well dispersed in the aluminum powders. As reported in our earlier publications, high energy ball milling has proven its effectiveness in dispersing the CNTs within the aluminum powders [7,18,19]. This is, however, accompanied by severe work hardening of the resulting composite powders as well as some coarsening in the powder particle size which necessitates controlling the amount of process control agent added. Such factors could result in processing challenges. Our current investigation addresses such effects by conducting preliminary research on the consolidation of these mixed powders in order to develop a better understanding of the impact of such differences on the feasibility of producing dual matrix composites as well as to identify conditions that lead to poor results so that they can be avoided.

Experimental Procedures

Al (99.7 % pure, - 200 mesh, Aluminum Powder Company Ltd., UK) and multi-wall carbon nanotubes (MWCNTs) (approximately 140 nm average diameter and 7 um in length, supplied by the MER corporation, USA), were used in the present study. 5 wt.% CNT was used to prepare the composite powders with the balance being Al. The mixture was milled under argon in stainless steel jars using stainless steel milling balls (ball-to-powder ratio, BPR = 5:1) at 400 rpm for 30 min. The amount of methanol added as a process control agent (PCA) to reduce the particle welding was limited to 50uL. To prevent heat build-up during the milling process which would result in the formation of undesirable reaction products at the interface of CNTs and aluminum, the ball milling process was interrupted every 10 min. Milled composite powders were then added to equal weight of unmilled aluminum powder (milled for 5 min at 200 rpm in order to break off the oxide layer) and were turbula mixed for 30 min to produce dual matrix powders. Those were then compacted at 475 MPa. Hot extrusion of the compact was conducted at 500°C using an extrusion ratio of 4:1. Dog-bone tensile test samples were machined and super-finished out of the extradâtes, and tested to fracture in order to generate a stress-strain curve.

Scanning electron microscopy (SEM) was used to determine the powder morphology following the milling process as well as investigating the fracture surface following the tension test in order

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to evaluate the interface strength. For microstructural characterization, extruded specimens were sectioned along the cross-section, ground and polished to 1 micron finish.

Results and Discussion

Initial efforts were focused on evaluating the particle size and morphology following the milling process as well as the dispersion of the CNTs. As can be seen from Figure 1, the particle size is not uniform; two distinct particle sizes are observed (18 urn and 75 urn). The presence of 5 wt% CNT favoured particle fracturing and thus a decrease in the particle size of the milled composite powders, as noted in our previous publication [19]. It is also noted that dispersed CNTs are present on the surface of the particles (Figure 1 b). Ideally, CNTs should be embedded within the particles as otherwise they could sustain damage under the impact of the milling media. Additionally, CNTs residing on the particle surface could interfere with subsequent particle bonding during the consolidation process. The addition of the PCA was necessary because trials in which the PCA was eliminated resulted in particle sticking to both the jars and the milling balls and thus low powder yield. Upon mixing the milled powders with equal amounts of unmilled aluminum powders the resulting morphology observed is presented in Figure 2.

Figure 1. SEM micrograph of (a) milled Al-5wt%CNT composite particles and (b) a higher magnification image of one of the particles showing several CNTs.

Figure 2. Dual powders (50% Al - 50% (Al-5wt% CNT)) mixed in the turbula mixer for 30 min.

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A representative stress-strain diagram of a dual matrix extruded sample is presented in Figure 3. A maximum strength of 165.5 MPa and a maximum strain of 1.3% (measured using a clip-on extensometer) represent the best results obtained. Such strength value which is 38% higher than the average value obtained for pure unmilled aluminum samples (typically 120 MPa), prepared under the same conditions, confirmed the reinforcing effect of the composite particles. However, this was accompanied with the undesirable effect of a severe loss in the ductility of the samples.

Figure 3. Stress strain diagram of a dual matrix (50% Al - 50% (Al-5wt% CNT)) sample.

Evaluation of the fracture surface by SEM revealed clearly (Figure 4 a) the difference in behavior between the soft unreinforced aluminum which developed large dimples as opposed to the heavily strain hardened milled composite particles which don't seem to have deformed as much. Figure 4 b is a higher magnification image which focuses on the interface between two particles and shows porosity along the interface due to the difference in deformation behavior between the two particles which appear to easily debond. The lack of a strong bond at the interface could also be attributed to the presence of some CNTs on the surface, as noted earlier, as well as other contaminants or oxide films. It is believed that the weak interface is responsible for inefficient load transfer from the outer soft aluminum matrix to the inner harder particles which lead to the observed unsatisfactory mechanical behavior.

Figure 4. Fracture surfaces of a dual matrix extruded sample showing (a) the two distinct regions (b) the interface between an unreinforced particle and a strain-hardened composite one.

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Figure 5. Microstructure of dual matrix composites showing the distribution of the CNT-A1 composite particles (dark) within the aluminum outer region. The inset shows details of the

CNT-A1 composite particles showing pores as well as oxide particles.

Figure 5 shows the microstructure of polished and etched samples. The inset focuses on a composite particle and reveals many pores surrounding the CNTs. Such poor bond would also contribute to the overall inadequate mechanical behavior.

Recent investigations attempting to overcome the previously explained shortcomings addressed the problem of CNTs residing on the particle surface as well as particle coarsening and non-uniform size by milling the powders for 1 hr instead of 30 min using a BPR of 5:1, as well as the use of higher purity PCA. Careful handling of the milled powders (which were kept under an argon atmosphere throughout the process) was also exercised in order to minimize any oxidation of the surface-active powders following milling. SEM observations (not shown) confirmed that the modified milling conditions lead to CNTs being embedded within the particles as well as a more uniform particle size. The milled composite powders were then mixed with equal weight of unmilled powders by further milling for an additional 1 hour. The final powder morphology is shown in Figure 6 and shows the more uniform particle size. Homogenization time before extrusion was extended to 1 hr instead of 30 min to ensure better consolidation.

Figure 6. Dual powders (50% Al - 50% (Al-5wt% CNT)) milled for an additional 1 hr.

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Figure 7. (a) Extruded dual matrix sample (b) Broken tensile test sample.

Figure 8. SEM micrograph of the fracture surface of a dual matrix extruded sample.

Preliminary results show a noticeable increase in the formability of the samples. The extruded dual matrix composites were free of extrusion defects at the surface (Figure 7 a), as opposed to the earlier samples in which surface cracks in the form of bamboo defects were frequently observed. In addition, noticeable enhanced ductility with consistent failure in the gauge length was observed (Figure 7 b). The initial results of the tension samples show much improved strength reaching 365 MPa as well as enhanced ductility reaching 6.7%. Fracture surface investigations (Figure 8) reveal overall ductility and less noticeable difference in the dimple size and depth between the two regions as opposed to the earlier samples in which ductility was only observed in the unmilled regions. This research is still the subject of ongoing work, but the results are encouraging and it is expected that this novel microstructural design would result in composites with reasonable ductility in addition to the enhanced mechanical properties. Samples with other CNT contents as well as different ratios of unmilled versus composite powders will also be prepared and evaluated.

Conclusions

A number of recommendations can be made regarding the dual matrix samples:

1. The dual matrix design has the potential to produce samples of high strength and good ductility. 2. Controlling the milling conditions so as to yield fine uniform particles as well as embedded CNTs is desirable.

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3. Ensuring a clean metallurgical interface by careful handling of the milled powders to minimize oxidation is essential.

Further work is underway to compare the dual matrix composites to single matrix ones processed under the same conditions. Preparation of fracture samples that allow us to characterize the fracture behavior is also in progress.

Acknowledgments

The authors wish to acknowledge the financial support of the US-Egypt Joint Science and Technology fund (grant number MAN11 -011-007) and the National Science Foundation (Office of International Science and Engineering) under grant number 0710869, the Yousef Jameel Science and Technology Research Center (YJSTRC) and the office of the Associate Provost for Research Administration at the American University in Cairo, Egypt.

References

1. T.K. Kyung, C. Seung, H.H. Seong, H.H. Soon, "Microstructures and tensile behavior of carbon nanotube reinforced Cu matrix nanocomposites", Materials Science and Engineering A, 430 (2006), 27-33. 2. C.S. Goh, J. Wei, L.C. Lee, M. Gupta, "Effect of fabrication techniques on the properties of Carbon nanotubes reinforced magnesium composites", Solid State Phenomena, 3 (2006) 179-182. 3. W. Daoush, B-K Lim, C.B. Mo, D.H. Nam, S.H. Hong, "Electrical and mechanical properties of carbon nanotube reinforced copper nanocomposites fabricated by electroless deposition process", Materials Science and Engineering A, 513-514 (2009) 247-253. 4. D. Poirier, C. Goujon, R. Gauvin, R. Drew, "Fabrication of aluminum matrix composites reinforced with carbon nanotubes", Advanced Materials and Processes, 163 (2005) 45. 5. H. Kwon, M. Estili, K. Takagi, T. Miyakazi, A. Kawasaki, "Combination of hot extrusion and spark plasma sintering for producing carbon nanotube reinforced aluminum matrix composites", Carbon, 47 (2009) 570-577. 6. R.George, K.T.Kashyap, R.Rahul, S.Yamdigni, "Strengthening in carbon

nanotubes/aluminum composites", Scripta Materialia, 53 (2005) 159-1163. 7. A.M.K. Esawi, K. Morsi, A. Sayed, A.A. Gawad, P. Borah, "Fabrication and properties of dispersed carbon nanotube-aluminum composites", Materials Science and Engineering A, 508 (2009) 167-173. 8. K. Morsi, A.M.K. Esawi, P. Borah, S. Lanka, A. Sayed, " Characterization and spark plasma sintering of mechanically milled Aluminum- Carbon Nanotube (CNT) composite powders", Journal of Composite Materials, 44 (2010) 1991-2003. 9. S.R. Bakshi, V. Singh, K. Balani, D.G. McCartney, S. Seal and A. Agarwal, " Carbon nanotube reinforced aluminium composite coating via cold spraying", Surface & Coating Technology, 202 (2008) 5162-5169. 10. K. Morsi, A.M.K. Esawi, S. Lanka, A. Sayed, M. Taher, " Spark plasma extrusion (SPE) of ball-milled aluminum and carbon nanotube reinforced aluminum composite powders ", Composites Part A: Applied Science and Manufacturing, 41 (2010) 322-326. 11. H.J. Choi, G.B. Kwon, G.Y.Lee, and D.H. Bae, "Reinforcement with carbon nanotubes in aluminium matrix composites", Scripta Materialia, 59 (2008) 360-363.

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12. T. Tokunaga, K. Kaneko and Z. Horita, "Production of aluminium-matrix carbon nanotube composite using high pressure torsion", Materials Science and Engineering A, 490 (2008) 300-304. 13. A.M.K. Esawi and M.A. El Borady, "Carbon Nanotube-Reinforced Aluminium Strips", Composites Science and Technology, 68 (2008) 486-492. 14. V.C. Nardone, J.R. Strife, K.M. Prewo, "Microstructurally toughened particulate-reinforced aluminum matrix composites", Metal. Mater. Trans. 22A (1991) 171. 15. Y. Chen, H.M. Wang, "Microstructure and wear resistance of laser melted TiC reinforced nickel aluminide dual-phase matrix in situ composite", Intermetallics, 14 (2006) 325-331. 16. X. Deng, B.R. Patterson, K.K. Chawla, M.C. Koopman, C. Mackin, Z. Fang, G. Lockwood, A. Griffo, "Microstructure/hardness relationship in a dual composite", Journal of Materials Science Letters, 21 (2002) 707-709. 17. K. Morsi, A.M.K. Esawi, P. Borah, S. Lanka, A. Sayed, M. Taher, "Properties of Single and Dual Matrix Aluminium-Carbon Nanotube Composites Processed via Spark Plasma Extrusion (SPE)", Materials Science & Engineering A, A 527 (2010) 5686-5690. 18. A.M.K. Esawi and K.Morsi," Dispersion of carbon nanotubes (CNT) in aluminium powder", Composites Part A: Applied Science and Manufacturing, 38 (2007) 646-650. 19. K. Morsi and A. Esawi, "Effect of Mechanical Alloying Time and Carbon Nanotube (CNT) content on the evolution of Aluminum (Al)-CNT Composite Powders", Journal of Material Science, 42 (2007) 4954-4959.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INFLUENCE OF HIGH PRESSURE TORSION ON THE

CONSOLIDATION BEHAVIOR AND MECHANICAL PROPERTIES

OF AA6061-SiCP COMPOSITES POWDERS

H.G. Salem1*, W.H. El-Garaihy2 and El-sayed M.A. Rassoul 3

'Mechanical Engineering Department, Youssef Jamil Science and Technology Research Center (YJSTRC), American University in Cairo, Egypt.

2Mechanical Engineering Department, Faculty of Engineering, Suez Canal University, Egypt. 'Production and Mechanical Design Department, Faculty of Engineering, El_Mansoura University, Egypt.

Keywords: high pressure torsion, powder consolidation, AA6061-SiC composite

Abstract

Disks of AA6061 reinforced with 0 and 15%SiC were processed via a combination of hot compaction (HC) of the mixed powders followed by high pressure torsion (HPT). HPT was conducted using two deformation modes; continuous rotation in a forward direction (monotonie (m-HPT)) and reverse rotation (cyclic (c-HPT)) over a number of revolutions under 1-GPa pressure at room temperature. Hardness profiles of sections cut parallel and perpendicular to load direction revealed that one revolution produced nonuniform distribution of strain-hardening, where highest hardness was displayed at the disk edges and lowest towards the centers. Increasing number of revolutions via m-HPT enhanced the hardness uniformity across the discs. Although, c-HPT produced equiaxed structure with high hardness, it was unable to eliminate the nonuniform distribution of properties across the discs even after 4 revolutions. HPT processing increased the Hv of HC disks by 280 and 200% for AA6061 with and without SiC reinforcement, respectively.

Introduction

6xxx series Al alloys are widely used in aerospace engineering for their good corrosion resistance and low cost, along with good formability and weldability. It is now realized that a better understanding of strengthening and flow properties of the Al alloy by manipulating the grain refinement processes can enable more effective exploitation in various applications [1], The most interesting materials commercially utilize SiC, AI2O3, Si3N4, TiC, and Ti3Al4

particles incorporated into the aluminum matrix by a variety of processes including powder metallurgy (PM) [2-3]. The primary property improvements being utilized are increased elastic modulus, enhanced wear resistance, it possess higher specific strength, specific stiffness, and wear resistance, improved high cycle fatigue performance [4-5].

It is realized that, one of the efficient methods to induce ultrafine grain (UFG) structure in materials is deforming it to large strains below recrystallization temperature without intermediate thermal treatment. Conventional metal forming processes such as forging, extrusion, drawing, or rolling do have limited steam intensity that is capable only of producing a grain size > 2 um. On the other hand, severe plastic deformation (SPD) techniques such as equal-channel angular pressing (ECAP), and high-pressure torsion (HPT), are capable of producing UFG solids with grain size ~ 0.2 urn [6]. Moreover, SPD has the ability of reducing porosity during deformation compared to conventional processes [7], Accordingly, a combination of PM and HPT was proposed and investigated in the current research to produce bearing materials with enhanced mechanical properties through retention and refinement of the consolidated ultrafine initial composite and monolithic powders.

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Materials, Processing, and Characterization

Micron-powders of AA6061 alloy with an average size of 30 urn were employed as the matrix, and SiC paniculate with mean size of 2 urn was adopted as reinforcement. The as-received AA6061 powders were supplied by the aluminum powder Company limited (APC) and the SiC powders were supplied by the American Elements Company. The as-received powder of SiC had irregular shape with a size ranging between l-5um. To fabricate the AA6061-15%v SiC metal matrix composite, the required fractions of AA6061 and SiC powders were mixed under vacuum in Argon atmosphere in a glove box followed by mixing in a turbula mixer operating at a rotation speed of 96 rpm. Both AA6061, AA6061-SiC composite powders were subjected to single sided uniaxial hot compaction (HC) into cylindrical sample 10 mm in diameter and 9.7 mm in height. The hot compaction process was carried out at compaction pressure of 525 MPa and temperature of 400°C, for 30 min.

The HC samples were subjected to HPT as a secondary consolidation processing step. Processing via HPT was conducted at room temperature using an apparatus described previously [8-9] with two massive dies; each die was machined to have a central depression of 4.4 mm deep and 10 mm in diameter, located centrally on its surface. Torsional straining was achieved by rotation of the lower die at a constant speed of 1 rpm. The upper surface of each disk was marked immediately after HPT and prior to any microstructural analysis and hardness measurements. HPT was conducted under pressure of P=lGPa. HPT was conducted using two deformation modes; continuous rotation in a forward direction or monotonie HPT (m-HPT) through total number of up to N = 4 revolutions, this deformation mode was denoted (route A) throughout the manuscript. The second deformation mode is cyclic deformation or cyclic HPT (c-HPT). For c-HPT, the first direction of straining (clockwise or counter-clockwise) was designated F, the reverse direction was designated R. The c-HPT process was carried out either through (1F+1R), this was designated 2 revolutions of route B (2-B) or though (1F+1R+1F+1R) which were designated 4 revolutions of route B (4-B).

Following HPT, the disks polished to a mirror-like finish. Vickers hardness (Hv) was measured on the disk sections cut parallel and perpendicular to the loading direction. Hv profiles were recorded on the disk surfaces separated by spacing's of 0.5mm [8, 10]. Color-coded contour maps depicting the variation of the local hardness across the surface of each sample were produced. Microstructural evolution and consolidation behavior of the disks before and after HPT were characterized by field emission scanning electron microscopy (FESEM) of ground polished and etched samples using "keller" solution.

Results and Discussion

Hardness Variation as a function of HPT Routes

Dependence of the hardness results on the distance r from the center of the disk is shown in Figure 1 for the m-PHT and c-HPT routes. The hardness contours of the HC samples were not displayed because they were essentially homogeneous across each disk with an average of Hv ~ 60 and 80 for AA6061 and AA6061- SiC, respectively. After 1 revolution of route A (1-A), the hardness values of AA6061 increased by a factor of~ 2 over the whole disk but at the same time the hardness distribution exhibited property inhomogeneous distribution as shown in Fig. la. This inhomogeneity was manifested in the central region where the hardness values displayed were lower within a diameter of ~ 6mm. knowing that AA6061 alloy has low stacking fault energy, recovery occurs at a slower rate, which produces a steady sate increase in the rate of hardening through the initial stages of processing. Lower recovery rate resulted in a high initial strain hardening that was associated with friction at the disc-die

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walls interface compared to lower hardness in the central region [10, 11]. At the central regions, hardness inhomogeneity was gradually lost with additional straining via 2-A as shown in Figure lb. Hv-values remain unchanged around the periphery of the disk while the lower hardness central regions decrease in diameter to 4 and 2mm with increasing the number of revolutions via 2-A and 3-A, respectively (Figure lc). The lower hardness central regions narrowed down to a limited zone with average Hv=119 after 4-A while the Hv values remained reasonably constant over the entire disk with an average of Hv=:125 as shown in Figure Id. This is consistent with the results produced by other researchers [12-13] and with the basic concepts of the strain gradient plasticity modeling that was developed to explain the evolution of homogeneity with increasing numbers of revolutions when processing in HPT [14]. Thus, the final increase in hardness is by a factor very close to double the values measured for the as hot compact discs of AA6061.

From Figure 1 (e-f) corresponding to c-HPT routes 2-B and 4-B, it was apparent that the hardness values were slightly higher at the center of the disk (Hv~102 and 119, respectively), which extended over a larger diameters compared to m-HPT sample processed via 2-A and 4-A routes (Hv~ 98 and 118, respectively). Similar to route A, route B processing produced inhomogeneous central core region with a diameter of- 4.6 and -1.3 mm for the 2 and 4 revolutions of route B, respectively. However, despite the differences in the hardness values at the central regions of the disks compared to route A, the values of Hv were essentially identical at the peripheries of the disks, which agreed with [15], For AA6061- SiC composite discs, intense plastic straining via HPT didn't eliminate the introduced inhomogeneity in the consolidated HC discs. This was manifested by the clear decrease in Hv values at the center of the disc after one revolution via route A (Figure lg). SiC addition resulted in an overall increase in hardness from 89-to-148 and from 122-to-200 Hv at the center and peripheries of the discs, respectively post one-revolution via route A. Increasing the number of revolutions resulted in an increase in Hv both at the center (184Hv) and peripheries (220Hv) up to 4 revolutions via route A, which was significantly higher than that displayed by the monolithic AA6061 discs as shown in Figure l(b-d). 4-revolutions via route A produced almost homogeneous distribution of hardness [10], while the maximum hardness measured did not change with increasing strain. It is worth noting that SiC addition reduced the positive influence of increasing the number of m-HPT revolutions on the increased property uniformity across the disc cross section. In other words, SiC reinforcement, although enhanced the overall discs hardness compared to the monolithic ones, promoted the inhomogeneity in hardness distribution even after 4-revolution. This could be attributed to the increased friction of the hard-SiC particles at the disc-die wall interfaces which increased strain hardening at the peripheries compared to the discs centers.

Figure l(k, 1), shows the measured Hv values for AA6061-SiC processed discs via c-HPT 2 and 4 revolution of route B. Although, almost similar Hv values were displayed within the central zones of the discs processed via 2 and 4 revolutions routes A and B, it was clear that the diameter of the central zones was larger for route B discs, which indicates lower degree of uniform distribution of properties [16]. Figure 2(a, b) show the Hv distribution on sections cut longitudinally for AA6061 and AA6061-SiC discs processed via 4 revolutions route A , respectively. It is clear that higher values of Hv were displayed at the peripheries compared to the central zone. Figure 2 also revealed more homogenous distribution of Hv for AA6061 compared to AA6061-SiC. AA 6061-SiC processed composite discs displayed 4-hardness zones, two of which represent severe strain hardening of ~224 Hv centered along the height of the disc and extending towards the center [9]. The remaining 2-hardness-zones represent lower strain hardening of ~195 Hv centered along the top and bottom surfaces of the discs.

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Figure 1. Hv profiles and color-coded contour maps showing Hv distribution across the surfaces of (a-f) AA6061, and (j-1) AA6061-SiC processed via: (a, g) 1-A, (b, h) 2-A, (c, i) 3-

A, (d, j) 4-A, (e, k) 2-B, (f, 1) 4-B.

Microstructural Evolution

The microstructure of the as-HC samples of AA6061 is shown in Figure 3a. The microstructure revealed the formation of almost equiaxed grains with average grain size of 35um with some porosity scattered at triple points along the boundaries of the consolidated powders. The SEM micrograph of AA6061-SiC composite in Figure 3b revealed the segregation of the SiC particles along grain boundary and at triple junctions (pointed at by arrows). HPT processing via 2A images show the consolidated powders particles elongation in the direction of shear stress for the AA6061 alloy discs with (Figure 3d) and without (Figure 3c) SiC additions.

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Figure 2. Hv profile on a plane eut parallel to the loading direction after HPT processing via 4 revolutions of route A for (a) AA6061, and (b) AA6061-SiC.

The angle of orientation of the structure of each condition reflects the degree of straining that the structure went through, it represent the losses in energy in terms of rigid body rotation and densification of the structure [17, 18]. SiC-addition to the Aluminum matrices induced higher degree of strain hardening of the structure, manifested by the observed excess elongation (higher aspect ratio) of the sheared grains compared to the monolithic AA6061 post 2-revolutions via route A. It is suggested that fragmentation of the SiC particles along the boundaries and their impingement on the relatively soft Al-matrices (Figure 3d) is an indication of the induced high strain hardening when composite discs are subject to SPD.

Figure 3. SEM micrographs of (a, c) AA6061 as-HC and (b, d) AA6061-SiC composite before and after m-HPT via 2-A route, arrows point at SiC particles

Microstructural evolution of the disk surfaces at the peripheries post HPT processing via 1 and 4 revolutions of route A, and 4 revolution of route B are shown in Figure 4 (a-c) for AA6061 and (d-f) for AA6061- SiC samples. Influence of the deformation amount and mode on the shape and size of the subgrain structure developed during deformation is clearly depicted in the images displayed. It is clear that increasing the amount of induced strain from

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l-to-4A resulted in sever elongation of the subgrain structure as shown in Figures 4 (a, d) compared to (b, e), respectively. 4-revolutions c-HPT produced equiaxed subgrains compared to elongated ones for the monotonie m-HPT which shows different orientation within the each individual grain based on the inclination of the slip planes of the Al-matrices. SiC-addition to the Al-matrices (Figure d-f) resulted in induced higher strained structure, which was depicted by the observed severely deformed subgrains post 1 and 4 revolutions via route A (Figure 4 d, e). SiC particles are represented by arrows in Figure 4 d-f.

Figure 4. SEM micrographs of (a-c) AA6061 and (d-f) AA6061-SiC processed via (a, d) 1 -A, (b, e) 4-A, and (c, f) 4-B, arrows point at SiC particles.

Grain, subgrain and subcell sizes were revealed via high resolution field emission scanning electron microscope as shown in Fig 5. In order to explain the influence of HPT processing on the consolidation behavior and the mechanical properties of the processed alloy and its composite, the following structural features will be discussed. Intensity of deformation was observed via the alignment of the AA6061 consolidated particles (grains) in the direction of shear, which increased with increasing the number of revolutions via m-HPT route 2-A and 4-A as shown in Figure 5a compared to b and c, respectively. Due to the SPD induced with increasing number of revolutions, the consolidated powders equiaxed subgrains were heavily elongated but rather in different directions following the orientation of slip planes as shown in Figure 5b and c. Even within the same grain, several orientations of the subgrains revealed deformation on more than one slip plane of the aluminum matrix indicative of multiple slip, which was manifested in Figure 5c. This could provide an indication of the evolution of substructure of medium-to-high angle boundaries. Further investigation with TEM is necessary to validate this observation.

Increasing the amount of intense plastic straining via m-HPT up to 4-revolutions resulted in an obvious refinement not only of the grains but also the subgrain structures as shown in Figure 5a compared to b and c, respectively. The grain and subgrain sizes of the as-HC structures were reduced from 35 and 3.2um to 20 and 1.9um, and 20 and 1.3um for the monolithic and composite discs, respectively. Accordingly, the measured high Hv post 2 and 4 revolutions via routes A and B; can be explained by the refinement of the internal structure of the HC AA6061 with and without reinforcement post HPT processing. SiC-particles

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distribution along the grain boundaries of the consolidated sheared grains increased the resistance to deformation around the ceramic hard particles and hence intensified the strain hardening, which was associated with dynamic recovery. A combination of strain hardening and dynamic recovery might have resulted in the observed increased refinement associated with the increased misorientation on the subgrain level [19]. This in-turn influenced positively the hardness level measured post HPT processing.

Figure 5.SEM micrograph for HPT processed samples via m-HPT (a)l-A,(b) 2-A and (c) 4-A

The present results show that the use of c-HPT produces a region of high hardness at the centers of the disks compared to m-HPT, which could be attributed to the finer grain structure produced compared to m-HPT. Hv results showed that the areas of microstructural inhomogeneity in the central regions of the disks extend to higher areas in c-HPT than m-HPT, it deceased with increasing numbers of revolutions in m-HPT whereas the inhomogeneous central regions were retained to higher numbers of revolutions when using c-HPT which is agreed with [16]. Such behavior of the structure evolution is known to be a consequence of the Bauschinger effect [20], which supports the suggested formation of HAGBs in the case of cyclic straining, is significantly retarded compared to the monotonie counterpart [20]. Accordingly, the c-HPT may be more effective in producing an equiaxed microstructure [18].

Conclusions

Disks of AA6061 reinforced with 0 and 15%SiC were successfully processed via combination of HC of the mixed powders followed by HPT. HPT was conducted using m-HPT and c-HPT up to 4-revolutions. The following conclusions can be drawn:

1. One revolution of HPT results in an inhomogeneous distribution of strain hardening, with higher values at the disk peripheries and decreased value towards the center. Uniformity of hardness distribution increased with increasing the number of revolutions.

2. Hv of the HC disks increased by 200 and 280% post HPT processing for AA6061 consolidated powders in the monolithic and composite conditions, respectively.

3. SiC-addition to AA6061 matrices results in decreasing the degree of uniform distribution of hardness across the disc longitudinal and transverse sections.

4. The variation of Hv values throughout the sample decreased with increasing the amount of shear strain, especially for m-HPT routes.

5. The area of microstructure inhomogeneity displayed at the centre regions of the disks extends over wider areas in c-HPT compared to m-HPT.

Acknowledgements

The authors are grateful to Youssef Jamil Science and Technology Research Center (YJSTRC), and the Mechanical Engineering Department, The American University in Cairo-

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Egypt, for financial support and continuous assistance of their lab engineers and research assistances.

References

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2. R. Fernândezl and G.G. Doncel, "Additivity of reinforcing mechanisms during creep of metal matrix composites: Role of the microstructure and the processing route", J. of Alloys and Compounds, 475(2009), 202-206. 3. G. Wan, "Bulk Al/SiC nanocomposite prepared by ball milling and hot pressing method", Trans. Nonferrous Met. Soc, 16 (2006), 398-401. 4. C.S. Shin and J.C. Huang, " Effect of temper, specimen orientation and test temperature on the tensile and fatigue properties of SiC particles reinforced PM 6061 Al alloy", International Journal of Fatigue, 32 (2010), 1573-1581. 5. T.W. Kim," Determination of densification behavior of Al-SiC metal matrix composites during consolidation processes", Mat. Science and Engineering, A 483 (2008), 648-651. 6. T.G. Langdon, "Processing of aluminum alloys by severe plastic deformation", Mat. Science Forum, Vols. 519-521 (2006), 45-54. 7. S. Seal et al., "Challenges and advances in nanocomposite processing techniques", Mat. Science and Engineering, 54 (2006), 121-285. 8. A. P. Zhilyaev and T. G. Langdon," Using high-pressure torsion for metal processing: Fundamentals and applications", Progress in Materials Science, 53 (2008), 893-979. 9. T.G. Langdon, G. Sakai, K. Nakamura, and Z. Horita, "Developing HPT for use with bulk samples", Mat. Science and Engineering, A, 406 (2005), 286-273. 10. T.G. Langdon, and C. Xu, "Three-dimensional representations of hardness distributions after processing by high-pressure torsion", Mat. Sei. and Engineering, A 503 (2009), 71-74. 11. T.G. Langdon, C. Xu, and Z. Horita "The evolution of homogeneity in an aluminum alloy processed using high-pressure torsion", Acta Materialia, 56 (2008), 5168-5176. 12. A. Loucif, R. B. Figueiredo.T. Baudin,F. Brisset, T. G. Langdon, "Microstructural evolution in an Al-6061 alloy processed by high-pressure torsion", Mat. Science and Engineering, A 527 (2010), 4864-4869. 13 A.P. Zhilyaev, K. Oh-ishi, T.G. Langdon, "Microstructure evolution in commercial purity aluminum during HPT", Mat. Science and Engineering A 410-411 (2005), 277-280. 14. Y. Estrin et al., "Strain gradient plasticity modelling of high-pressure torsion", J. of the Mechanics and Physics of Solids, 56 (2008), 1186-1202. 15. T.G. Langdon, C. Xu, and Z. Horita, "The evolution of homogeneity in processing by HPT", Acta Materialia, 55(2007), 203-212. 16. D. Orlov, Y. Tokada, M. Umemoto, and N. Tsuji, "Role of strain reversal in grain refinement by severe plastic deformation", Mat. Sei. and Eng., A 499 (2009), 427-433. 17. J. Zhang, N. Gao, M. J. Starink, "Al-Mg-Cu based alloys and pure Al processed by high pressure torsion: The influence of alloying additions on strengthening", Materials Science and Engineering, A 527 (2010), 3472-3479. 18. M. Kawasaki, B. Ahn, and T. G. Langdon," Significance of strain reversals in a two-phase alloy processed by high-pressure torsion", Mat. Sei. and Eng., A 527 (2010), 7008-701 19. R.Z. Valiev et al.," Nanostructures and Microhardness in Al and Al-Mg Alloys Subjected to SPD", Materials Science Forum, 604-605 (2009), 179-185. 20. T.G. Langdon, "The significance of strain reversals during processing by high-pressure torsion", Mat. Science and Engineering, A 498 (2008), 341-348.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

LASER POWDER DEPOSITION OF ALMGBi4-TIB2 ULTRA-HARD COATINGS ON TITANIUM, AND STEEL SUBSTRATES

J. Fuerst, M. Carter, J. Sears,

South Dakota School of Mines and Technology 501 E St. Joseph Street, Rapid City, South Dakota, 57701, USA

BAM, Boride, Wear Resistant, Lubricious

Abstract

BAM is an ultra-hard ceramic compound created, at Ames Laboratory in 1999, by combining Aluminum Magnesium Boride (AlMgBu) and Titanium Diboride (T1B2). The high hardness, low coefficient of friction, and relative high temperature chemical inertness of BAM make it an ideal material for wear resistant coatings. Current BAM coatings are produced by Physical Vapor Deposition (PVD) or magnetron sputtering and are capable of only thin films. Using LASER Powder Deposition (LPD), high hardness (1300 HV) BAM coatings of 1 mm thickness, with increased fracture toughness were produced on Ti-6A1-4V and mild steel substrates. Initial testing has demonstrated LASER Powder Deposition to be a capable means of manufacturing large, highly wear resistant surfaces.

Introduction

AlMgBi4 is one of a number of high hardness, boron rich compounds with a crystal structure influenced predominantly by the B12 icosahedra. The B12 icosahedron is treated as a single species, induced into orthorhombic staking by the presence of interstitial aluminum, magnesium, and boron [1], The high hardness of the material is believed to be primarily due to the highly ordered covalent bonding of the boron in the B12 icosahedra [2]. Further evaluation shows incomplete occupancy of all of the potential B12 icosahedra interstitial sites with approximately 25% and 22% vacancy by the aluminum and magnesium respectively [3, 4], The presence of the interstitial aluminum, magnesium, and boron alters the bond length between the B12 icosahedra producing a low degree of symmetry in crystal structure [1], Figure 1 shows the unit cell, comprised of 64 atoms, noting the random orientation of B12 icosahedra and incomplete occupation of all lattice vacancies [5].

Figure 1. Unit cell of AlMgBn showing low symmetry of the B12 icosahedra [6]

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The effect of the low symmetry of the crystal structure is that X-Ray diffraction of physical vapor deposition (PVD) and pulsed LASER deposition thin films of AlMgBi4/TiB2 (BAM) showed the resulting layer to be amorphous. AlMgBi4 peaks could not be identified at room temperature [5, 6],

Despite the low crystalliniry of the material, which does not normally indicate high hardness; testing of pure, nanocrystalline AlMgB14 demonstrated a microhardness of 32-25 GPa [7]. The addition of titanium diboride, TiB2, increased the microhardness to 35-46 GPa [7], Wear testing of BAM shows the material is self lubricating with a coefficient of friction of 0.08. The lubricity of BAM is due to the exposed boron reacting with atmospheric or environmental water and hydrogen to produce a thin film of boric acid, B(OH)3, at the surface of the material [8].

One experimental method used to create an engineered surface is LASER powder deposition. A LASER is used to melt a small volume of substrate material and a stream of fine powder is sprayed into the molten substrate using an inert carrier gas shown in Figure 2 below [9]. The thermal energy in the pool is what melts the sprayed powder. Because the path plan of the LASER control program keeps the LASER in motion across the surface of the material, a very small volume of the substrate is molten at any point in time. Only the material that is directly under the beam is momentarily molten. Because of the small volume of the melt pool, compared to the volume of the bulk substrate, the deposition solidifies very rapidly. The dimensions and mechanical properties of the deposition and the depth of the heat affected zone (HAZ) are dependent on the LASER energy, LASER spot size, substrate adsorption coefficient, deposition path plan, powder feed rate, and particle size.

Figure 2. Illustration of LASER deposition process [10],

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Experimental Procedures

Deposition Procedure:

All LASER powder deposition was performed using the M-LAM (Micro Laser Additive Manufacturing) VDK 3000 LASER Deposition System at the Additive Manufacturing Laboratory (AML) of the South Dakota School of Mines and Technology, Rapid City, South Dakota. A 1000 watt Nd:YAG LASER is contained within the system and is piped to the head using a 600 urn fiber. The LASER and robotic gantry are contained in a glove box under argon atmosphere.

The initial sample of BAM received was an agglomerate of TiB2 and AlMgBi4. The agglomerate size was sieved to -140 mesh. Initial testing revealed that the BAM material would not feed through the Mark XV/RD powder feeder. A second batch of powder was then prepared as a blend of BAM with 20% CP Ti. This blend was found to feed through the powder feeder, but only at high feed rates.

The argon flow rate for powder was 3.5L/min, the purge argon flow rate was 30L/min, and standoff distance between the deposition head and the substrate was 12.5 mm. However, LASER power, travel speed, and powder flow rate were all adjusted to optimize deposition.

A second sample of BAM powder was received, produced by Ceramatec (Salt Lake City, Utah) as an agglomerate of TiB2 and AlMgBu hot pressed into billet and crushed in a pulverizing process. The agglomerate size was sieved to -80 mesh. Testing revealed that the BAM material would feed through a Mark XV/RD powder feeder at most powder flow rates.

The argon flow rate for powder (3.5L/min), purge argon flow rate (30L/min), and standoff distance (12.5 mm) had not been adjusted from the initial depositions. However, LASER power, travel speed, and powder flow rate were all adjusted to optimize deposition. Samples were then deposited for wear testing and metallographic evaluation.

Hardness Testing and Metallography Procedure:

Vickers microhardness testing was performed using a Buehler (Lake Bluff, Illinois) Micromet 4 Micro Hardness Tester with a diamond indenter and an applied load of 300 g.

Selected samples were chosen for metallographic examination. Samples were cut on an Allied High Tech Products, Inc. (Rancho Dominguez, California) TechCut 5™ Precision Sectioning Machine with an Allied High Tech Products, Inc. aluminum oxide wafering blade. The hardness of the deposition prevented complete sectioning of the sample. The samples were placed in the machine so to cut through the substrate material. The deposition was then broken by means of a chisel. Samples were mounted using a LECO Corporation (St. Joseph, Michigan) PR-25 Mounting Press in a LECO Corporation black bakélite powder then polished using Allied High Tech Products Dia-Grid Diamond Disks.

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Wear testing was performed on an RPM and Associates, Inc. (Rapid City, South Dakota) Model NJ-1630 ASTM G65 Dry Abrasion Tester to the ASTM G65 specifications.

Results and Discussions

Figure 3A below shows the cross section of a BAM/20 wt% CP-Ti deposition on a 1045 steel substrate. Metallography shows that the sample had a mixed phase of CP- Ti and BAM while there is a distinct region at the surface that is exclusively BAM. Vickers microhardness testing was performed on the surface and main bulk of the deposition. Average Vickers microhardness was measured at 1345 HV in the regions of the deposition that, metallographically, appear to be free of CP-Ti. In regions of the deposition containing CP-Ti Average Vickers microhardness was measured at 879 HV.

Figure 3B below the cross section of a BAM deposition on 1045 steel substrate with a lower LASER energy. It can be seen that there is not as much segregation between the Ti phase and the BAM and full metallurgical bonding between the deposition and the substrate. The bulk of the deposition and the surface of the deposition appear to be homogeneous. Vickers microhardness testing of the full thickness of the cross section of the material confirmed that the sample was homogenous with an average Vickers microhardness 1213 HV.

Figure 3. a) Deposition of BAM with 20 wt% CP-Ti on 1045 steel, b) Deposition of BAM with 20 wt% CP-Ti on 1045 steel deposited with reduced LASER energy. The dark acicular particles in the deposition have been identified by Energy-dispersive X-ray spectroscopy to contain boron, so are assumed to be BAM. The lighter interstitial material has been identified as Titanium.

Figure 5A shows the cross section of a BAM/20 wt% CP-Ti deposition on a Ti-6Al-4Vsubstrate. The CP-Ti and BAM is homogenous through the entire deposition. Vickers microhardness testing of the full thickness of the cross section of the material confirmed that the sample was homogenous with an average Vickers microhardness 1157 HV. The decrease in hardness of the deposition of the BAM onto Ti-6A1-4V is theorized to have been caused by increased titanium dissolution into the deposition, diluting the concentration of AlMgBu. Figure 5B shows the

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cross section of a pure BAM deposition on 1045 steel. The sample has good dispersion of TiB2 and AlMgBi4 through the deposition. There is a well defined region of mixing of the BAM and the substrate material. Vickers hardness was performed on this sample and the bulk hardness was found to be 1415 HV.

Figure 5. A) Deposition of BAM with 20 wt% CP-Ti on Ti-6A1-4V. B) Deposition of pure BAM on 1045 steel.

Figure 6 shows the cross section of a pure BAM deposition on a Ti-6Al-4Vsubstrate with a LASER energy of 600 W. There is some visible desolution of the titanium substrate into the BAM of the deposition. This desolution decreases through the thickness of the deposition with the highest concentration of titanium adjacent to the substrate and a distinct region of nearly pure BAM at the surface. Vickers hardness tests were done at the surface and main bulk of the deposition. It was found that at the surface where there appears to be no Ti phase in the deposition the Vickers hardness was 1400 HV. Through the bulk of the part that did contain a Ti phase the Vickers hardness was found to be 500 HV.

Figure 6. Deposition of pure BAM on Ti-6A1-4V deposited at 600 W.

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X-Ray diffraction was performed on a few of the samples consisting of pure BAM deposited on a 1045 steel substrate in order to determine the deposition microstructure and composition (Figure 7). XRD showed the presence of iron boride at the surface of the deposition. This indicates that a small amount of dissolution of the BAM occurred within the steel melt pool and that it is believed that some of the interstitial boron was consumed to produce iron boride. Thermodynamic evaluation with the modeling software Thermo-Calc (Thermo-Calc Software, Inc., McMurray, Pennsylvania) indicates that this is a feasible mechanism of action. The XRD pattern indicates the AlMgBn predominately maintained a low symmetry microstructure required for high hardness.

Figure 7. XRD pattern for BAM deposited on steel showing formation of iron boride. Note the lack of aluminum or magnesium peaks due to the low symmetry of the material.

In order to increase surface hardness by reducing iron boride dissolution, the ASTM G65 Dry Abrasion Test wear coupons were produced using two layers of deposition. Results of the ASTM G65 Dry Abrasion Test were an average wear coupon volume loss of 2.574%. The ASTM G65 Dry Abrasion Test coupons can be seen below in Figures 8A and 8B as faced for testing and after wear testing respectively. Vickers microhardness testing of like samples produced with two layers of deposition had a measured hardness of 2221 HV.

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Table I. Percent Volume Loss for ASTM G65 Dry Abrasion Test

Samole 1 2 3 4

Initial mass (e) 184.152 159.774 176.078 147.164

Final Mass le) 184.143 159.764 176.069 147.156

Volume Loss % 2.574 2.860 2.574 2.288

Average: 2.574

Figure 8. ASTM G65 Dry Abrasion Test coupons A) after facing and B) after testing

Conclusion

Initial testing has shown that LASER powder deposition is a feasible means of producing highly wear resistant AlMgBi4/TiB2 BAM coatings on steel and titanium alloys. Continued parameter development needs to be performed, however, to achieve improved results. ASTM G65 Dry Abrasion Testing of deposited BAM demonstrated a wear resistance approximately three times higher than hard chrome plating and higher than NT-60 nickel tungsten carbide. Further testing will need to be done to enhance the mechanical properties of the deposition and reduce dissolution of the substrate material into the BAM as much as possible.

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References

1. Yongbin Lee and B.N. Harmon, "First Principles of Calculation of Elastic Properties of AlMgBi4" (Ames Laboratory and Department of Physics and Astronomy, Iowa State University, Ames, Iowa, August 22, 2002).

2. J. E. Lowther, "Symmetric Structure of Ultra Hard Materials" Journal of the American Ceramic Society, 85(1) (2002), 55-58.

3. Yongbin Lee, "First Principles Calculations for X-Ray Resonant Spectra and Elastic Properties" (Ph.D. Dissertation, Iowa State University, Ames Iowa, 2004).

4. R. Chattopadhyay, Surface Wear: Analysis, Treatment, and Prevention (Materials Park, OH: ASM International, 2001) 59.

5. Richard Bodkin, "A Synthesis and Study of AlMgBn" (Ph.D. Dissertation, University of the Witwatersrand, Johannesburg, South Africa, 2008).

6. Jason Curtis Britson, "Pulsed Laser Depostion of AlMgBu Thin Films" (Master's Thesis, Iowa State University, Ames Iowa, 2008).

7. Cook, B. A., J. L. Harringa, T. L. Lewis, and A. M. Russell. "A New Class of Ultra-Hard Materials Based on AlMgB14." Scripta Materalia 42 (2000), 597-602

8. Brice A. Cook, Joel L. Harringa, James Anderegg, Alan M. Ressell, Jun Qu, Peter J. Blau, Clifton Higdon, and Alaa A. Almoursi; "Analysis of Wear Mechanism in Low Friction AlMgB,4-TiB2 Coatings" Surface and Coatings Technology, 205 (2010) 2296-2301.

9. B. Vamsi Krishna, Weichang Xue, Susmita Bose, and Amit Bandyopadhyay; "Engineered Porous Metals for Implants" Journal of the Materials, Metals, and Minerals Society 60(5) (2008), 45-48.

10. James W. Sears; "Additive Manufacturing and Repair: How do Lasers Fit In?" 29* International Congress on Applications of LASERS & Electro-Optics. Hilton in the Walt Disney World Resort. Orlando, FL. October 29, 2007

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MECHANICAL PROPERTIES OF SPARK PLASMA SINTERED ZrC-SiC COMPOSITES

Sumbule Sagdic, Ipek Akin, Filiz Sahin, Onuralp Yücel, Gultekin Goller

Istanbul Technical University, Metallurgical and Materials Engineering Department, 34469 Maslak, Istanbul, Turkey

Keywords: ZrC; SiC; Microstructure; Spark plasma sintering; Mechanical properties

ABSTRACT

Zirconium carbide (ZrC) is an important structural ceramic due to its high melting temperature, excellent chemical resistance, high electrical conductivity and good mechanical properties. This combination of properties makes it potentially useful in application of cutting tools, high temperature crucibles and thermal protection components. However, the nature of brittleness and lack of damage tolerance is one of the most crucial problems in their applications. In this study, ZrC-SiC composites with different compositions were prepared by using spark plasma sintering (SPS) technique. Samples containing 10 and 20 vol% SiC were sintered at 1850°C for 300 s under a pressure of 40 MPa. Densities of the composites were determined by the Archimedes' method. Fully dense ZrC-SiC composites with a relative density of 99% were obtained. Mechanical properties and microstructural behavior of the composites were also investigated.

■ .Introduction

Zirconium carbide (ZrC) is a transition metal carbide which is characterised by high hardness (25 GPa), high melting temperature (~3420°C), high electrical conductivity (~106 S/m) and relatively high strength (E=390 GPa, G = 172 GPa). These properties give it the potential to be a useful engineering ceramic. However, the use of zirconium carbide in engineering applications has been limited by the lack of a fully developed, commercially viable sintering process. As an ultra-refractory compound, very high temperatures and pressure-assisted techniques are required to achieve dense bodies, due to its highly covalent bonding character, and low self diffusion coefficients. One promising approach for improving the densification is to use of field-assisted techniques such as spark plasma sintering (SPS). This technique employs a pulsed DC current to activate and improve sintering kinetics. Higher densities, refined microstructures, clean grain boundaries and elimination of surface impurities have been reported, which, in turn, result in an overall improvement in the materials' performance [1-3].

ZrC is a potential candidate for a variety of high-temperature structural applications such as jet engine parts, hypersonic vehicles, cutting tools, furnace elements, crucibles, etc. The realization of this potential is hindered, however, by two major drawbacks, namely the low fracture toughness of ZrC and the very high temperature required for its sintering because of covalent bonding [1], Also, it is difficult to obtain fully dense ZrC ceramics because of the rapid grains growth at high temperature leads to entrapped pores in the grains. A composite approach has

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been successfully adopted to overcome the low values of fracture tougness. Improvement in fracture toughness can be achieved by making composite through introduction of weak interfaces in to the material which act to deflect propagating cracks [2]. Effect of several sintering aids have been studied in order to lower the sintering temperature of ZrC based ceramics [4]. ZrC-SiC composites with different sintering additives (MoSi2, graphite) have been prepared by several researchers using spark plasma sintering [3], and hot pressing [1] techniques. However, there is no open literature on the spark plasma sintered ZrC-SiC composites without sintering aid.

2. Experimental Procedure

ZrC (Grade B, H.C. Starck Corp., an average particle size of 1-4 um) and a-SiC (Grade UF-10, H.C. Starck Corp., an average particle size of 1 um) powders were used as starting materials. Throughout the text, samples were designated by their component volume fractions. The raw materials were weighed in appropriate quantities, ball milled in ethanol for 24 h and then dried. A graphite die 50 mm in inner diameter was filled with the mixture, followed by sintering using SPS apparatus (SPS-7.40 MK-VII, SPS Syntex Inc.) at 1850°C for 300 s with a heating rate of 1.7°C/s in vacuum. A graphitic sheet was placed between the punches and the powder, and between the die and the powder for easy removal and better conductivity. A uniaxial pressure of 40 MPa and a pulsed direct current (12 ms/on, 2 ms/off) were applied during the entire SPS process. The pressure was released during cooling for all samples.

An optical pyrometer, focused on a small hole at the surface of the graphite die, was used to measure and adjust the temperature. The current was controlled manually. Linear shrinkage of the specimens during SPS process was continuously monitored by displacement of the punch rods. The sintered specimens were in the form of pellets 50 mm in diameter and 5 mm thick and characterized after sand-blasted in order to remove the graphitic sheet.

The crystalline phases were identified by X-ray diffractometry (XRD; MiniFlex, Rigaku Corp.) in the 28 range of 30-80° with CuKa radiation. The bulk densities of the specimens were determined by the Archimedes' method and converted to the relative densities according to the equation 1:

DR = ( D B / D T ) * 1 0 0 (1)

where DR is the relative density, DB is the bulk density (Mg/m3) of the sintered sample and DT is the theoretical density (Mg/m3) of the composite. The value of theoretical density depends on the composition and calculated using theoretical densities and volume fractions of ZrC (6.7 Mg/m3) and a-SiC (3.2 Mg/m3). The fracture surfaces of the specimens were observed by scanning electron microscopy (FESEM; JSM7000F, JEOL Ltd.). Vickers hardness (Hv) was measured under loads of 9.8 N and fracture toughness (Kic) was evaluated by microhardness tester (VHMOT, Leica Corp.), under load of 19.6 N from the half-length of a crack formed around the indentations by using the following Eq. (2), where P is load (N), Ec is Young's modulus (GPa) of composites calculated assuming a mixture rule, H is the Vickers microhardness (GPa), k is a constant and c is half of the average crack length:

Klc=k*(Ec/H)"2*(P/c15) (2)

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Eq. (2) was derived from Anstis et al. for median cracks. For a full dense composite, Young's modulus of the composite can be described by a rule of mixture as follows.

Ec=IEi*fi (3)

where Ej is young's modulus of ith constituent phase, and f* is the volume fraction of rth constituent phase, and n is total number of constituent phases. With the values of Ei=390 GPa (ZrC), E2=450 GPa (a-SiC), Young's modulus of composites in Eq. (3) yields 396 GPa and 402 GPa for the 90 ZrC-10SiC, 80ZrC-20SiC composites respectively. These theoretical Young's modulus values for dense structures were multiplied by the relative densities of the composites and calculated values of Ec were used in the measurements of fracture toughness. The average value of the 20 measurements for each sample was used for the evaluations of hardness and toughness.

3. Results and Discussion

3.1 Relative Density and Crystalline Phase

The theoretical densities of the composites were calculated according to the rule of mixtures assuming no impurities and no reactions during processing. In order to determine the relative density, the bulk density was divided by the theoretical density. The relative densities of the spark plasma sintered samples with varying addition of SiC are listed in Table 1. The relative density of ZrC increased from 95.5% to 99% with the addition of 10 vol% silicon carbide (SiC). The highest relative density, 99.1%, was achieved with the addition of 20 vol% silicon carbide (SiC) for the 80ZrC-20SiC composite.

Table 1. Compositions (vol%), relative densities, hardness and fracture toughness of ZrC and ZrC-SiC composites consolidated by SPS.

Composition (vol%) 100 ZrC

90ZrC-10SiC 80ZrC-20SiC

Relative density (%) 95.5 99.0 99.1

Vickers hardness (GPa)

17.7 ±0.2 18.2 ±0.3 20.0 ±0.3

Fracture toughness (MPa-mr/2) 3.3 ±0.2 4.3 ±0.2 4.8 ±0.3

Figure 1. shows the XRD pattern of ZrC-SiC composites containing 0, 10, 20 vol % SiC sintered at 1850°C for 300 s. Characteristic peaks of ZrC ( JCPDS:65-8704) and a-SiC (JCPDS:49-1428) were identified and no chemical reaction was detected between phases at 1850°C.

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Figure 1. The XRD pattern of lOOZrC, 90ZrC-10SiC, 80ZrC-20SiC composite sintered atl850°C for 300 s.

3,2 Microstucture

Microstructures of the fracture surfaces of ZrC-SiC composites containing 0, 10, 20 % volume SiC sintered at 1850 °C for 300s are shown in Fig. 2. The fracture surface of the ZrC particles showed features of cleavage.

The distribution of grains was homogeneous for all samples.The fracture surfaces of SiC containing ZrC-SiC composites indicated a dense microstructure and the morphology of the

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grains was mainly equiaxed. The grain size of ZrC and SiC were in the range of 3-6 (im and 2-4 um, respectively.

From the microstructure of the fracture surfaces of the sintered samples, it can be seen that the SiC (black phase) is homogeneous throughout the samples. There was a reduction tendency for grain growth in ZrC with increasing SiC amount. For the composite containing 10 vol% SiC (Fig. 2(b)), most SiC particles were distributed at the ZrC/ZrC grain boundaries and a small amount of SiC particles were entrapped in the ZrC matrix. In this case, the amount of SiC is comparatively low and die grain growtii of ZrC is only partially restained. At 20 SiC vol% nearly all SiC particles were dispersed at ZrC/ZrC grain boundaries and surrounded ZrC grains in the ZrC-20 vol% SiC composite in Fig. 2(c), which demonstrates that the SiC particles can effectively inhibit the grain growth of ZrC.

Figure 2. SEM micrographs of fracture surfaces of the 100 ZrC (a), 90ZrC-10SiC (b), 80ZrC-20SiC (c).

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3.3 Hardness ans fracture toughness measurements

Table 1 summarizes the effect of SiC adition on the Vickers hardness and fracture toughness of ZrC-SiC composites at load of 9.8N and 19.6 N, respectively. The hardness of the composites increased with increasing SiC content.

In order to determine the toughening mechanisms, interaction between propagating crack and microstructure was analyzed. The fracture surface of the ZrC particles exhibited features of cleavage as shown in Fig. 2. The cleavage of ZrC particles implied that when the propagating crack met a ZrC particle, the crack most frequently crossed through the ZrC rather than propagated around the particle and this is consistent with the observation on the crack path of ZrC-SiC composite containing 20 vol% SiC as shown in Fig. 3. The crack propagated along the ZrC grains and showed trans-granular tendency. In the SiC phases, the crack propagated along the grain boundaries and deflected at an angle (shown by D in figure 3) and according to the crack deflection model [5,6], the energy consumption increased. The crack deflecting effect of SiC could be responsible for the higher fracture toughness of SiC containing composites.

Figure 3. SEM image of the crack propagations of the polished surface of ZrC-SiC composites containing 20 vol% SiC.

Ma et al. have studied relative density, fracture toughness and Vicker's hardness of the ZrC-SiC-Cg composites prepeared by hot press and Sciti et al. have investigated the physical and mechanical properties of the spark plasma sintered ZrC with an additive of MoSi2. If the results given in Table 1 are compared with the above studies [1,3], where the combination of relative

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density, hardness and fracture toughness were achieved as 99.1%, 20 GPa and 4.8 MPa-m"2, respectively, it is obvious that the results obtained in this study are superior.

4. Conclusions

ZrC and ZrC-SiC composites containing 10 and 20 vol% SiC were prepared by spark plasma sintering at 1850°C for 300 s.The results of this study showed that fully dense ZrC-SiC composites with a relative density of 99.1% were obtained at 1850 °C. Vickers hardness of ZrC-SiC composites increased with increasing SiC content, and composite containing 80ZrC-20SiC sintered at 1850°C for 300 s had the highest value of 20 GPa. Fracture toughnesss of ZrC-SiC composites increased with increasing SiC content, and composite containing 80ZrC-20SiC sintered at 1850°C for 300 s had a highest value of 4.8 MPa-m"2.

Acknowledgements

The authors thank H.Dincer for his contribution in SPS studies, H. H. Sezer and T. T. Alpak for SEM studies.

References

1. B. Ma et al., "Fabrication of hot-pressed ZrC-based composites," Proceedings of the Institution of Mechanical Engineers Part G-Journal of Aerospace Engineering, 223 (G8) (2009), 1153-1157.

2. B., Ma, W., Han, "Thermal shock resistance of ZrC matrix ceramics," International Journal of Refractory Metals and Hard Materials, 28 (2) (2010), 187-190.

3. D., Sciti, S., Guicciardi, M., Nygren, "Spark plasma sintering and mechanical behaviour of ZrC-based composites," Scripta Materialina, 59 (6) (2008), 638-641.

4. L. Zhao et al., "Pressureless sintering of ZrC- based ceramics by enhancing powder sinterability," International Journal of Refractory Metals and Hard Materials, 29 (4) (2011), 516-521.

5. K.T. Faber, A.G. Evans, "Crack deflection processes-I.Theory", Ada. Metall. 31 (1983) 565-576

6. K.T. Faber, A.G. Evans, "Crack deflection processes-II.Experiment", Ada. Metall. 31 (1983)577-584

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INTENSE PULSED LIGHT SINTERING TECHNIQUE FOR NANOMATERIALS

H. A. Colorado1'3'*, S. R. Dhage4, J. M. Yang' and H. T. Hahn2'5

'Materials Science and Engineering, University of California, Los Angeles, USA 2Mechanical and Aerospace Engineering, University of California, Los Angeles, USA

Universidad de Antioquia, Mechanical Engineering. Medellin-Colombia 4International Advanced Research Center for Powder Metallurgy & New Materials, India

5School of Mechanical & Aerospace Engineering, Seoul National University, Korea

" Corresponding author. Tel: +1-217-778-3728 Email: hcoloradolopera@ucla.edu

Abstract

The Intense Pulsed Light (IPL) is an emerging technique that uses optical light for thermal processing. Current applications include sintering and processing of nanomaterials for printed electronics. IPL can sinter nanoparticles in just milliseconds. Such a short reaction time prevents oxidation of the elements and second phase generation. Another valuable benefit of IPL sintering is that materials can be irradiated without damaging the glass or low temperature substrates such as flexible polymers. The solid state diffusion of nanoparticles in a very short reaction time of few milliseconds is the great benefit of IPL system over the conventional smtering methods. It is believed that if the particles have very high surface area-to-mass ratio, very little light is needed to sinter the nano-particles. Melting and recrystallization of particles to larger grains without structural deformation and phase transformation are possible because of very short reaction time. Thus, IPL can be a very promissory technique for ultra rapid processing of nanomaterials and thin films. Several case studies are briefly presented.

Keywords: Intense pulsed light, Photonic sintering, Nanomaterials, Printed Electronics.

1. Introduction

Nano and electronic materials typically require sintering techniques that provide just the right amount of low energy to produce local melting, in order to avoid degradation of the properties. On the other hand, metallic nano materials when exposed to high temperatures show more oxidation than micro or bigger configurations due to the high surface area in nano materials. Intense Pulsed Light (IPL) sintering gives solutions to the problems presented above. Figure la shows a classification for different sintering techniques, divided in three major fields: pressureless, pressure-assisted and electromagnetic field-assisted. IPL and photonic sintering fall into the first category that includes sintering under air, in a pressureless process. Sintering under pressure (allows obtaining dense ceramics at low temperature with low grain growth) includes the most commonly used methods (hot pressing and hot isostatic pressing). Other methods include electromagnetic fields as for instance microwave sintering and spark plasma sintering. In the latter, pressure is applied during the process.

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Sintering has three main steps [1,2]: a) Bridging or necking of particles, which is a formation of a connection region between the

grains. It is activated by diffusion mechanisms, evaporation condensation, plastic deformation, etc. and finishes when bridges have grown near 50% of the grain radii. Typically in this step mere is an increase of about 15% in density.

b) Elimination of residual cavities or pores, which magnitude depends from surface energy. Since now cavities are interconnected by bridges, interconnected pores appear like a network, which decrease with time and lead into a compactness of near 90%.

c) Disappearance of almost all pores, giving a near full dense material. This is the step where grain growth can be important.

Figure 1. a) different sintering techniques, b) sintering steps [1]

The processing temperature in solid state sintering is typically slightly above two-thirds the temperature of the melting point. Sintering has many variables which can be divided in the type of powder (morphological such as shape or size and chemical as purity) and the processing (such as temperature, pressure, environment and sintering duration).

The IPL [3-7] uses a light that typically has a spectrum of wavelengths from 200 to 1100 nm, which means that more of the used light for processing is in the visible range. In addition, because light can be focused, it can be intense enough to produce localized melting. Typically light is emitted in pulses as short as microseconds duration, which decrease the change to materials to be oxidized. IPL is a type of Photonic sintering at high intensities and short pulse duration.

This paper shows several case studies including applications in thin films and printed electronics, conductive inks and processing of nanomaterials. The main characteristics of the technique are also presented.

2. Pulsed light sintering technique

Details of the IPL system have been described previously [1]. The main parts in a typical IPL system are: a xenon flash lamp, an aluminum reflector, a power supply, capacitors, a simmer triggering pulse controller, and a light filter. The irradiation is generated by using arc plasma

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with the xenon flash lamp. Typical flash lamp emits a spectrum of light of wavelengths from 200 to 1100 nm. To distribute the high electrical current (about 1000 A) when the lamp is triggered, a charged super capacitor (40000 uF) can be used. All this occurs between 1 andlO ms.

For the conversion energy, approximately 50% of the electrical energy is converted into radiated optical power, in the wavelength we used: 200-1100 nm. As an example of a typical test, see as described later in Figure 3, the IPL system was configured to generate a pulse of two energy peaks of 50 J/cm"2. The duration of the pulse is between 3 to 4 ms each and separated by 20 ms to avoid heat saturation on the lamp. Every pulse corresponds to two shots. Thus 2, 5, 10, 20, 50 and 100 irradiation shots were performed over the samples.

2.1 Light source and flash lamps

Typical lamps produce light in a UV/Visible spectrum from 200nm to lOOOnm [8, 9], with special lamp spectral cutoffs wavelength: A: 370nm, B: 240nm and C: 190nm. IPL is generated using an arc plasma phenomenon in the xenon flash lamp [10]. Figure 2a shows a typical IPL lamp. Figures 2b and c show a representation of the light spectra with different cutoffs wavelength [8] and a typical spectral distribution for a flash lamp respectively. Figure 2d shows a representation of IPL sintering in printing electronics.

Figure 2. a) Typical IPL lamp and b) Light spectra with different cutoffs wavelength [8]; c) typical spectral distribution of the flash lamp between 200 and 80nm [10], d) representation of

IPL sintering in printing electronic

579

Once the gas is ionized, a spark is formed between the electrodes, allowing the capacitor current to conduct, which is the arc plasma phenomenon.

3. IPL applied for the preparation of printed electronics

In printed electronics functionally conductive inks and coatings contain nanoparticles that permit the use of low-cost substrates such as polymers and papers. Nowadays silver, gold and copper nanoparticle inks can be used in inkjets and screen printers in a sintering process that can sinter or anneal particles at low temperatures, typically below 160°C. This is possible because IPL techniques use high energy peak pulses, typically with few milliseconds duration, which quickly heat the inks and not the substrates. During the process, the high energy quickly removes the solvent where the particles are, leaving only the metal particles mat are being treated while the substrate is not affected. Thus, it is possible to print on substrates sensitive to high temperatures such as flexible printed circuit boards [8], Since the duration of the high energy peak pulses is too short, it allows metal particle inks to be sintered without forming oxide layers that affect the conductivity and other electronic properties. However, it has some challenges to solve, some common to pressure-less sintering, for example the presence of residual pores which increases the resistivity of the manufactured film. It has been pointed out for Kim et al [4] for copper nanoink that surface of the obtained films has a different morphology when is compared to the bulk, which has being associated to a critical penetration length of light where pores can be minimized.

IPL has been successfully proven in several substrates such as Photo paper, Copier paper, Label paper, Polyethylene terephthalate (PET), Polyimide, Polyethylene film, Polycarbonate, Polystyrene and other polymeric films. Inkjet, Gavure and Flexographic print methods have successfully used to apply the inks [9, 10]. In addition, sintering can be conducted at temperatures as low as 50°C with sintering times as low as 1ms. Table I summarizes parameters of results for IPL sintering applied to printed electronics.

3.1 Precious metals-based nanoinks

A considerable amount of on conductivee lines and the substrates for printed electronics has been conducted on using conductive polymers despite their conductivity and durability are much lower compared with metallic conductive lines. Typical materials such as poly(acetylene), poly(pyrrole), poly(thiophene), polyaniline, polythiophene, poly(p-phenylene sulfide), and poly(p-phenylene vinylene) are commonly used as conducting lines [11].

Gold and silver nanoparticles have been used for conductive ink fabrication due to its high conductivity and thermal stability. However, these particles are too expensive for mass production.

J. Kang et al. [12] has fabricated Ag nanoink with IPL by using Ag nanoparticles dispersed in diethylene glycol (DEG), see Table I. It was found that Although the power intensity decreased, the sintering quality of the silver nanoink pattern improved as the number of consecutive pulses was varied from one to two to three consecutive pulses, keeping the total energy density constant

580

at 50 J/cm2. It was also found that IPL irradiation of 50 J/cm2 increased the temperature of the printed silver nanoink pattern by 500°C.

3.2 Copper nanoinks

Although metal- based inks made of Au and Ag are widely used due to their excellent conductivity and easy manufacturing at low temperatures under conventional processes, they are too expensive which limits the implementation. Recent developments on copper (Cu) nanoinks [13] and IPL gave solution to the problems of using expensive Ag and Au particles. IPL, by applying short pulses of intense light irradiation, enable the manufacturing of printed electronics by a millisecond sintering without oxidation. Cu nanoparticles are cheaper than precious metals nanoparticles, but are easily oxidized under ambient conditions, then, a protective over-layer may then be applied after the process to enhance long-term performance. Also, the melting temperature of copper is relatively higher than that of silver.

Kim et al. [4] found that when light intensity was higher than 40J/cm2, agglomerate particles appeared. On the other hand, when light intensity was lower than 30J/cm2, large cracks appeared. Thus, a light intensity of 50J/cm2 (with a lamp 14mm away from the substrate) was used to produce IPL sintered copper nanoink without damages on the polymer substrates. The pulse duration was fixed to 2ms. The sintered film had 5uI2 cm electrical resistivity, which is one third (17.2uTi cm) of the thermally sintered copper nanoink reported by Park et al. [13]. The instantaneous temperature was estimated to be 600-700 instantly.

Han et al. [14] conducted a systematic study on the effect of IPL sintering of Cu nanoinks. Pulse length, the number of pulses, and the total energy were investigated. In order to avoid possible damage on the substrate for one energetic light shot, multi-pulses where applied while maintaining the total energy over the threshold level. Optimum sintering conditions were obtained at pulse energy 32 J/cm2, number of pulses of more than 4 and pulse length less than 3ms. Other methods like reactive sintering for Cu nanoparticles with IPL has been recently developed [15]. In this technique, IPL simultaneously sinters and reduces metal oxide nanoparticles.

581

Table 1. Some reported results for IPL sintered metallic nanoinks for printed electronics

Sintered material/substrate

Ag nanoparticles (GFRPC laminate; bismaleimide triazine (BT) was used as the polymer matrix) [121 Cu nano ink/(glass fiber bismaleimide triazine, epoxy composite, polyimide film, PP film and PE film [41

Cu nano ink/Polyimide (PI) substrate [14]

Cu nano particle with CuÛ2 shell and functionalized poly (N-vinylpyrrolidone) (PVP)coatinR layer [151

Pulse length Imsl

14 2 3

2

1.5 1.5 1.5 3.3 2.4 1.5 12 6 4 3

1.5

2

Gap between pulses Imsl

5 5

12 12 12 12 11 10

84 42 28 12

-

Energy |J/cmz|/number of pulses

50/1 50/2 50/3

50/1

12/8 22/8 32/8 32/8 32/8 32/8 32/1 32/2 32/3 32/4 32/8

50

electrical resistivity lull ml

0.186 0.089 0.047

5

1.78 1.73

Distance lamp-target Imml

14

14

4

4

4

Particle size [nm|

20-40

5

30

30

30

30

4. IPL for processing nanomaterials

4.1 Sintering nanoparticles

Intense Pulsed Light (IPL) has been used for processing CdS nanoparticles [5] as shown in Figure 3a. It was found that the number of irradiation shots conducted with the IPL technique increased the crystallinity of the CdS and the formation of CdS nanopillars. A simple mechanism of ultra fast melting and cooling like quenching has been proposed to explain the phenomenon. The major advantage of intense the pulsed light from Xenon flash lamp (IPL) treatment was to improve the crystallinity without inducing phase transformation. Additionally, this transformation has been achieved in a very short time (a few milliseconds). This extreme treatment was conducted in such short annealing time that prevented the oxidation/phase transformation of CdS and secondary phase generation. Another benefit of IPL treatment was that the CdS nanoparticles were annealed without damaging the glass/flexible polymer support material (Figure 3).

582

Figure 3. a) Scheme of the intense pulsed light with the xenon flash system and its corresponding pulse-energy-time representation, b) IPL energy pulse representation and XRD results snowing

that crystallinity increases as number of shots increases [5]

5. Summary

The IPL technique for printed electronics by sintering nanoparticles in nano metallic inks and for processing semiconductor nanomaterials has been discussed. Benefits of IPL include sintering or annealing without damaging die glass/flexible support material, which is generally sensible to high temperatures; ultra fast irradiation shots allow the process being done in few milliseconds, which eliminates the oxidation problem of diermal sintering under air atmosphere; irradiation area is in the order of centimeters; and all these advantages lead into a inexpensive powerful technique, perfect for mass production.

The IPL is a technique with great promise and should be applied to a much wider range of materials for the processing of semiconductors and metallic powders for thin films and bulk applications.

Acknowledgements

The authors wish to thank to COLCIENCIAS from Colombia for the grant to Henry A. Colorado.

6. References

1. D. Bernache-Assollant. Chimie-Physique du frottage (Ed: D. Bernache-Assollant), Hermes Forceram Collection, Paris 1993, Ch. 7.

583

2. T. Hungria, J. Galy and A. Castro. Spark plasma as a useful technique to the nanostructuration of piezo-ferroelectric materials. Advanced Engineering Materials 2009, 11, No. 8. Pp 615-631.

3. J. W. Sears and M Carter. Low temperature sintering of printed nano-particulate materials for fabrication of passive electronics. IMAPS Advanced technology workshop on Integrated/Embedded passives. San Jose, CA, 2007.

4. H. Kim, S. R. Dhage, D. Shim, H. T. Hahn. Intense pulsed light sintering of copper nanoink for printed electronics, Appl. Phys. A (2009) 97:791-798.

5. H. A. Colorado, Sanjay Dhage and H. T. Hahn. Thermo chemical stability of CdS nanoparticles under Intense Pulsed Light and high temperature XRD. Materials Science and Engineering B 176 (2011) 1161-1168.

6. Saad Ahmed. Low temperature Photonic Sintering for Printed Electronics. Xenon, 2011. http://semiconwest.org/sites/semiconwest.org/files/3_Saad%20Ahmed_Xenon.pdf

7. Sanjay Dhage, H. A. Colorado and H. T. Hahn. Morphological variations in CdS nanocrystals without phase transformation. Nanoscale Research Letters 2011, 6:420.

8. XENON Corporation, Pulsed Light Sintering R&D System. Http//:www.xenoncorp.com 9. Novacentrix, Advanced curing for printed electronics. Http//:www.novacentrix.com 10. http://optoelectronics.perkinelmer.com/content/RelatedLinks/CAT_flash.pdf H.H. Sirringhaus, T. Kawase, R.H. Friend, T. Shimoda, M. High-Resolution Inkjet Printing of

All-Polymer Transistor Circuits. Inbasekaran, W. Wu, and E.P. Woo, Science 290, 2123-2126 (2000).

12. J.S. Kang, J. Ryu, H.S. Kim and H.T. Hanh. Sintering of Inkjet-Printed Silver Nanoparticles at Room Temperature Using Intense Pulsed Light. Journal of Electronics Materials, Vol. 40, No. 11,2011

13. B. K. Park, D. J. Kim, S. H. Jeong, J. H. Moon, J. S. Kim. Thin Solid Films 515, 7706-7711 (2007).

14. Won-Suk Han, Jae-Min Hong, Hak-Sung Kim and Yong-Won Song. Multi-pulsed white light sintering of printed Cu nanoinks. Nanotechnology 22 (2011) 395705 (6pp).

15. J. Ryu, H-S Kim and H. T. Hahn. Reactive sintering of copper nanoparticles using intense pulsed light for printed electronics. Journal of Electronic Materials, Vol. 40, No. 1, 2011.

584

TIMIS2012 141 s t Annual Meeting & Exhibition

Recent Developments in Biological, Electronic,

Functional and Structural Thin Films and Coatings

Edited by:

Nuggehalli Ravindra Jian Luo

Xing Yang (Mark) Liu Nancy Michael Roger Narayan

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DEPENDENCE OF TRIBOLOGY OF CARBIDE DERIVED CARBON FILMS ON HUMIDITY

TMS1, Marcin Tlustochowicz2

'TMS (The Minerals, Metals & Materials Society); 184 Thorn Hill Rd.; Warrendale, PA 15086-7514, USA

2CTLGroup; 5400 Old Orchard Road; Skokie, IL, 60544, USA

Keywords: Tungsten; carbide; carbon; friction; humidity; tribocorrosion; tribolayer; film; chlorination.

Abstract

Tribologically advantageous films of carbide derived carbon (CDC) have been successfully synthesized on binderless tungsten carbide by reacting it with chlorine at 1000°C. Some of the treated samples were later dechlorinated by an 800°C hydrogénation treatment. The results of detailed characterizations of the CDC films and sliding contact surfaces were correlated with the friction and wear behavior of the CDC films in various tribosystems, including CDC-steel, CDC-WC, and CDC-CDC and in two specific environments: moderately humid air and dry nitrogen, as well as in an environment of changing water vapor pressure. Friction coefficient values as low as 0.11 in moderately humid air and 0.03 in dry nitrogen were obtained. A model of tribological behavior of CDC has been proposed that takes into consideration the tribo-oxidation of counterface material, the capillary forces from adsorbed water vapor, the carbon-based tribofilm formation, and the lubrication effect of both chlorine and hydrogen.

Introduction

In the past fifteen years, a new technique has been developed that allows synthesis of carbon coatings by etching of non-carbon species from carbides in a halogen-containing gas mixture. The resulting carbide derived carbon (CDC) can be produced on various metal carbides at ambient pressure and moderately elevated temperatures [1, 2]. The CDC films can be grown on any desired shapes with their dimensions preserved, and the growth rate is very quick compared to other processes [1]. CDC has a unique microstructure, non-uniform and consisting of diamond structures close to the interface with the carbide, carbon onions and graphitic spheres further away from the interface and closed-shell carbon structures and planar graphite furthest from the interface [2].

The most widely researched carbide for synthesis and evaluation of CDC coatings is silicon carbide, SiC. This work is an attempt to examine the tribological performance of CDC synthesized on binderless tungsten carbide manufactured by a plasma sintering process [3] while the findings are compared to CDC products derived from silicon carbide.

On the other hand, still relatively little is understood about the complex tribological behavior of CDC films. Previous research on the subject [1, 2] concentrated on the adhesion aspect of friction and wear, neglecting the important processes of tribocorrosion, abrasion, capillary action of adsorbed liquids, and transfer layer (tribofilm) formation. In this work, a more comprehensive and thorough depiction of CDC tribology is provided, based on processing tungsten carbide and subsequent study of CDC films synthesized on it.

587

Methods

Tungsten carbide samples used in this study were supplied by Materials Modification, Inc., based in Fairfax, VA. They were manufactured using their plasma pressure compaction (I^C) method [3] from pure tungsten carbide powder. The larger samples were cut into square coupons of approximate size of 18 x 18x2 mm. The coupons were then polished with diamond pastes of 15 um, 9 um, 6 (im, 3 um, and 1 um particle size. CDC was synthesized on the coupons in a furnace with a horizontal fused silica tube with an inner diameter of 60 mm and a length of 180 cm. A 3.5% Ch-Ar gas mixture, chlorination temperature of 1000°C, and times of 1 h to 3 h were used, resulting in CDC films of thicknesses varying from 20 um to 50 urn, depending on the synthesis time. The chlorination reaction can be described as follows:

WC(s) + 2Cl2(g) -*WCl4(g) + C(s) (1)

In order to remove undesirable residual chlorine and improve the tribological performance of CDC synthesized on WC in certain environments, some samples were subjected to post synthesis 8-hour heat treatments at 800°C using a 5.0% H2-Ar gas mixture where residual chlorine was reacted with hydrogen and formed hydrogen chloride gas that could then vaporize and leave the system together with the Fb-Ar gas mixture flow. The hydrogénation reaction is as follows:

H2 + C12/C1" — HC1 (g) (2)

Treated samples were then analyzed using Raman spectroscopy and scanning electron microscopy (SEM) to determine the presence of carbide-derived carbon structures. Subsequently, CDC films were subjected to ball-on-disk testing using a CSEM tribometer and two types of counterbodies: M50 steel and WC-Co 9.55 mm diameter balls. Steel balls were specifically used to determine if residual chlorine remaining after treatments has a detrimental effect on tribological results for steel counterbodies. The friction coefficient test parameters were as follows: sliding velocity of 0.15 m/s, 1 N load (for comparison with previous studies), and room temperature of 21°C. Two general conditions were used: open laboratory air (13-36% relative humidity) and dry nitrogen (less than 0.5% relative humidity).

After 4-hour ball-on-disk tests were performed in air for the investigated tribopairs, wear track surfaces and sliding contact surfaces of the balls were analyzed using SEM with energy dispersive spectroscopy (EDS) and a MicroXAM optical microscope and MapVueEX v. 6.16 surface profilometry software. Elemental compositions of transfer layers between the rubbed surfaces were determined using EDS. Specifically, the wear track on untreated binderless WC after it was rubbed against M50 steel was examined, as well as the tribopairs of M50 steel ball -CDC, and WC-Co ball - CDC. At least 5 measurements were taken for each surface and then the average was calculated. Surface profilometry was also performed on M50 steel balls after static corrosion tests where the balls were in static contact with CDC surfaces.

Finally, an additional test using the CSEM tribometer was designed to study the dependence of the friction coefficient values of CDC on humidity, where the tribopair was a CDC ball (SiC based) against a CDC disk (WC based). After the ball-on-disk test was started in dry nitrogen and a steady-state low friction coefficient was achieved, air was introduced into the testing environment and water vapor pressure was controlled in four steps at 50 Pa (2% RH), 100 Pa (4% RH), 230 Pa (10% RH) and 700 Pa (30% RH). Then, the system was purged with dry

588

nitrogen to achieve 0% RH again while the friction coefficient was constantly measured. The remaining test conditions were the same as for the other counterbodies.

Results and Discussion

In a previously published paper [4], a marked improvement was reported in moderately-humid-air friction coefficient values of CDC after its hydrogen post-treatment when compared to as-synthesized CDC. For the WC counterbody, the friction coefficient drops from -0.21 to -0.11 and, for the steel counterbody, from -0.26 to -0.14. Chlorine levels in chlorinated samples, determined by EDS more recently, were observed at average levels of 0.22 at.% - 1.48 at.%. The hydrogen post-treated samples exhibited average chlorine levels of 0.04 at.% and 0.10 at.%, which shows that hydrogénation can help in removal of up to 95% of chlorine from treated WC samples, which in turn leads to lower friction coefficient values in moderately humid air. The removal of chlorine during hydrogénation was not complete most probably because it appears to be bonded to the CDC structure which will be discussed later.

Surface profilometry performed on M50 steel balls after static corrosion tests, where the balls were in static contact with as-synthesized and hydrogen post-treated CDC films for 5 days, showed that corrosion of counterbodies by the action of residual chlorine in CDC (forming HC1 on contact with moisture) was significant and that tribocorrosion plays an important role in increasing friction coefficient values in moderately humid air for as-synthesized CDC without hydrogen post-treatment while in absence of moisture in dry nitrogen atmosphere this role is greatly diminished.

The results of EDS elemental analysis of the wear tracks for M50 steel ball tribopairs are presented in Table I. While a layer of tribo-corrosion and abrasion products, containing mainly iron oxides, forms on untreated WC, no similar layer forms on the CDC film with iron oxides and carbon forming a layer on the steel ball. This shows that CDC films can improve frictional properties of WC tribosystems and prevent galling, or sticking of the machined steel material to tungsten carbide surface—an unwanted phenomenon concerning WC parts and tools. It is worth noting that the transfer layer on the steel ball contains small quantities of chlorine which is a good indicator that it is chemically bonded within the carbon structure of the formed film.

Table I. EDS elemental analysis of the wear tracks on untreated WC and CDC synthesized on WC and the wear scar on an M50 steel ball after ball-on-disk tests.

Surface Type Wear track on untreated WC after sliding against M50 ball Wear track on CDC synthesized from WC after sliding against M50 ball Wear scar on M50 ball after sliding against CDC

At% W

8.4

1.5

-

At% Fe

33.3

1.0

25.0

At% C

12.0

92.3

58.0

At% Cl

-

0.2

0.5

At% O

46.3

5.0

16.5

The EDS transfer layer investigation results for the WC-Co ball - CDC film tribopair, are presented in Table II. A transfer layer on the wear scar of the WC-Co ball was identified to consist of carbon and various oxides, although different than those on the steel ball. A high level of chlorine in the transfer layer on the WC-Co ball (the highest of all measured) might again indicate that Cl is chemically bonded within the graphitic film structure and that chlorine specifically travels from other areas of CDC to take advantage of some unsatisfied bonds within

589

graphite. Generally, the levels of chlorine were found to be higher in the wear tracks on the CDC films and suggest that chlorine might bind to help remove unsatisfied bonds in the rearranged graphitic structures that form during sliding which in turn lowers friction.

Table II. EDS elemental analysis of the wear track on CDC synthesized from WC and the wear scar on a WC-Co ball after ball-on-disk tests.

Surface Type Wear track on CDC synthesized from WC after sliding against WC-Co ball Wear scar on WC-Co ball after sliding against CDC

At% W

0.4

7.3

At% Co

-

3.3

At% C

94.0

63.6

At% Cl

1.0

2.7

At% O

4.6

23.1

Optical microscope examination and profilometry were performed on the steel balls that were used in the 4-hour ball-on-disk tests on both as-synthesized and hydrogen post-treated CDC surfaces. Sample three-dimensional visualization and optical micrographs are shown in Figure 1. The difference in the wear of the balls that slid against as-synthesized and hydrogen post-treated CDC is very noticeable. The ball rubbed against as-synthesized CDC was flattened considerably and very little transfer material from the CDC film formed on its wear scar. The ball used in the test against hydrogen post-treated CDC, on the other hand, showed almost no wear and exhibited strong formation of a transfer layer. These observations indicate that the presence of chlorine in CDC (in absence of hydrogen post-treatment) might help increase the wear of steel several times and at the same time it might prevent the formation and growth of the transfer layer on the surface of the steel ball.

There is a very significant improvement in friction coefficient values for both counterbody types when the ball-on-disk tests are performed in a dry nitrogen atmosphere (Figure 2). In dry nitrogen, as-synthesized (not hydrogen post-treated) CDC exhibits friction coefficient values as low as 0.028 against both M50 steel and WC after an initial transition/transfer layer formation period for both counterbody types that lasted about 4 hours. However, further hydrogen post-treatment does not appear to improve the friction coefficient results which remain similar to as-synthesized CDC values. Therefore, it can be conferred that, for applications in dry-nitrogen environments, no hydrogen post-treatment of CDC synthesized on binderless WC is necessary.

Optical microscope examination and profilometry shows no significant difference between steel ball wear against as-synthesized and hydrogen post-treated CDC surfaces (Figure 3). It appears that the corrosion action of chlorine on the counterbody material when in sliding contact with CDC in air might be greatly reduced in a dry nitrogen environment. Additionally, once the corrosion effect on frictional properties is diminished, formation of graphitic tribolayers on the counterbody surfaces becomes easier and might start to play a decisive role in obtaining low friction coefficient values.

Virtually identical frictional behavior in dry nitrogen of as-synthesized CDC films, containing up to 1.5 at.% of residual chlorine, and hydrogenated CDC films, where hydrogen replaces most of the chlorine, leads to the conclusion mat either bout hydrogen and chlorine play a similar role in the reduction of the strong unsatisfied bonds within the highly disoriented CDC structure, or they both do not.

590

Figure 1. Optical micrographs and 3D visualization of wear scars of steel balls slid in air against untreated (left) and dechlorinated (right) CDC films.

Figure 2. Friction coefficient vs. time for ball-on-disk tests in dry nitrogen on binderless WC chlorinated at 1000°C before and after hydrogen post-treatment.

591

Figure 3. Optical micrographs and 3D visualization of transfer layers on steel balls slid in dry nitrogen against untreated (left) and dechlorinated (right) CDC films.

Some evidence, that chlorine might indeed help reduce the dangling bonds on the surface of CDC and also in the CDC-based tribofilm on the counterbodies, comes from the EDS analyses showing higher chlorine concentrations in the transfer layer and in the wear tracks on CDC than overall in CDC films. If chlorine molecules were not satisfying these bonds they would probably not get transferred to the tribofilm but would get released in the process of its formation.

Additionally, another halogen gas, fluorine, has been successfully employed to reduce the dangling bonds in graphite [5] and DLC [6] so chlorine may similarly bind in CDC structures.

If, on the other hand, neither chlorine nor hydrogen play a role in satisfying the strong dangling carbon bonds on CDC surfaces, the explanation of how they are reduced would be difficult. Carroll at al. [7] suggested that carbon onions present in CDC might rearrange during sliding and help close off the edge planes in graphite structures within CDC but this idea seems far-fetched since it is obvious that hydrogen or chlorine molecules can move much easier within CDC than any large structures.

The results of the ball-on-disk test where water vapor pressure varied are presented in Figure 4 and indicate a very strong dependence of frictional properties of CDC on water vapor pressure with the biggest change between 0% RH and 2% RH.

592

Figure 4. Friction coefficient vs. time for a ball-on-disk test in changing water vapor pressure.

This phenomenon can be explained as follows: when CDC is initially rubbed against itself in dry nitrogen, a graphitized tribolayer starts forming with carbon rearranging in such a way that the interaction between the two surfaces is reduced only to weak types of bonding like van der Waals or n-it* interactions. In addition, the tribolayer that forms between the two surfaces that is very easy to shear which in turn provides extremely low friction coefficient values. Finally, the absence of any oxidation products reduces the possibility of entrapped matter between the two surfaces which would increase abrasion. When water vapor is introduced into the system, it is immediately re-adsorbed to the surface. The tenfold difference in friction coefficients (0.03 to 0.3) corresponds to the difference in the strength of van der Walls forces or JI-JT* interactions and hydrogen bonding of water molecules to carbon. Additionally, strong capillary forces come into play as soon as a layer of water appears on the rubbing surfaces.

Based on the findings presented above, a model of tribological behavior of CDC was created. While the three main components of friction in CDC tribosystems are abrasion, shearing and adhesion, mechanisms controlling corrosion and erosion are important, as well. The lowest friction coefficients in CDC are achieved by the formation of a transfer layer that works in two ways: shields the counterbody material from CDC (possibly containing chlorine), reducing or even eliminating its corrosion and erosion and subsequently abrasion, and serves to absorb shearing during sliding and dissipate some energy this way.

Additionally, the capillary forces cease to exist when the adsorbed water vapor is removed (in case of a dry nitrogen atmosphere) and the resulting unsatisfied bonds in the structures of CDC and tribofilm are most probably satisfied by either chlorine or hydrogen (in case of hydrogenated CDC) which in turn leaves only the weakest forces available for adhesion between the two rubbing surfaces.

The tribological behavior of CDC in humid air is governed mostly by the water vapor pressure but, in case of as-synthesized CDC (without dechlorination), the situation is additionally complicated by various corrosion processes caused by residual chlorine that include oxidation of

593

the counterbody which in turn promotes abrasion, increases the oxide content in the tribofilm, or might even prevent its formation.

Conclusions

The research presented in this paper has thoroughly demonstrated that CDC films with beneficial tribological properties can be successfully grown on binderless tungsten carbide manufactured through the plasma pressure compaction (P2C) method.

When CDC is used on WC, it reverses the process of galling by creating a transfer film on the steel counterbody so both WC and steel surfaces are protected.

Ultra-low friction coefficient values of 0.03 in a dry nitrogen environment have been confirmed for CDC grown on binderless WC when they were previously reported for hydrogenated CDC synthesized on other carbides. However, it was determined that no dechlorination of CDC (i.e. in the form of hydrogen post-treatment) is necessary to obtain this unique behavior.

These very low steady-state friction coefficients in dry nitrogen can be achieved thanks to the formation of a carbon-based tribofilm that both shields the counterbody from wear and corrosion while at the same time providing low interfacial shear. However, water vapor and the corrosive action of chlorine control the frictional behavior of CDC in humid air and can cancel the benefits of the tribofilm or sometimes even prevent it from forming.

References

1. Y. Gogotsi, I-D. Jeon, and M J. McNallan: "Carbon Coatings on Silicon Carbide by Reaction with Chlorine-Containing Cases", Journal of Materials Chemistry 1997, (7) 1841-1848.

2. A. Erdemir, A. Kovalchenko, M. J. McNallan, S. Welz, A. Lee, Y. Gogotsi, B.Carroll: "Effects of High-Temperature Hydrogénation Treatment on Sliding Friction and Wear Behavior of Carbide-Derived Carbon Films," Surface & Coatings Technology 2004, (188-189) 588-593.

3. T. S. Srivatsan, R. Woods, M. Petraroli, and T. S. Sudarshan: "An Investigation of the Influence of Powder Particle Size on Microstructure and Hardness of Bulk Samples of Tungsten Carbide," Powder Technology 2002, (122) 54-60.

4. M. Tlustochowicz, M. J. McNallan, A. Erdemir and T. S. Sudarshan: "Tribology of Carbide-Derived Caron Films Synthesized on Tungsten Carbide," Surface Modification Technologies XXII Proceedings, T. S. Sudarshand and P. Nylen eds., Valar Docs, 2008.

5. R. L. Fusaro, H. L. Sliney: "Graphite Fluoride (CFx)n - A New Solid Lubricant," ASLE Tranactions 19 70, 13 (1) 56-65.

6. C. Donnet, J. Fontaine, A. Grill, V. Patel, C. Jahnes, M. Belin: "Wear-Resistant Fluorinated Diamondlike Carbon Films," Surface and Coatings Technology 1997, (94-95) 531-536.

7. B. Carroll, Y. Gogotsi, A. Kovalchenko, A. Erdemir, M. J. McNallan: "Effect of Humidity on the Tribological Properties of Carbide-Derived Carbon (CDC) Films on Silicon Carbide," Tribology Letters 2003, 15(1) 51-55.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

STRUCTURAL AND OPTICAL PROPERTIES OF SILICON CARBONITRIDE THIN FILMS DEPOSITED BY REACTIVE DC

MAGNETRON SPUTTERING

Okan AGIRSEVEN", Tolga TAVSANOGLU*, Esra OZKAN ZAYIM**, Onuralp YUCEL*

'Istanbul Technical University, Metallurgical and Materials Engineering Department, 34469, Maslak, Istanbul, Turkey

agirseven@itu.edu.tr, tolgatav@itu.edu.tr, yucel@itu.edu.tr Istanbul Technical University, Faculty of Science and Letters, Engineering Physics

Department, 34469, Maslak, Istanbul, Turkey ozesra@itu.edu.tr

ABSTRACT

In this study, silicon carbonitride thin films of variable compositions were deposited on glass and AISI M2 high-speed steel substrates by reactive DC magnetron sputtering of high purity silicon target using CH4 and N2 as reactive gases. The composition of the coatings was modified by the change in the reactive gas flow ratios. Microstructural properties were investigated by cross-sectional SEM analyses. Spectrophotometer was used to measure the optical transmittance and reflectance of silicon carbonitride thin films over the spectral range from 280 to 1000 run. The optical constants and band gap values of the films were further evaluated with respect to the gas flow rate. The results of analyses and calculations provided the information about the relationship between the reactive gas flow rates, microstructure, optical constants and band gap values of silicon carbonitride films.

Keywords: silicon carbonitride, thin film, reactive magnetron sputtering, microstructure, band gap, optical constants

1. Introduction

In recent years, by the increase in the need for wear resistant coatings, silicon carbonitride films have gained significant attention in engineering applications for their superior mechanical, tribological and optical properties. Silicon carbonitride films show good strength, thermal stability, wear resistance and corrosion resistance. SiCN films are used as protective coatings because of their high oxidation resistance and thermal stability. They are also high performance electronic materials. SiCN ceramics are used in MEMS systems because of their chemical inertness and high hardness [1,2,3], SiCN is a wide band gap semiconductor [4]. Adjustable optical properties and band gap characteristics make SiCN thin films suitable for microelectronic and optoelectronic applications [2],

SiCN films can be deposited by a variety of techniques, such as high temperature chemical vapor deposition, radio frequency plasma enhanced chemical vapor deposition, pyrolysis of polymer precursors, electron cyclotron resonance PECVD, pulsed laser deposition, ion beam sputtering and reactive magnetron sputtering [2]. From the various possible choices, magnetron sputtering

595

appears to be very attractive, due to its relative simplicity, high attainable deposition rates and wide acceptance by industry [4].

2. Experimental

2.1. Coating

Silicon carbonitride thin films of variable compositions were deposited on glass substrates by reactive DC magnetron sputtering of a single crystal silicon target of 99.999% purity with a diameter of 150 mm and a thickness of 7 mm, using TSD 350 PCVD model hybrid coating machine (HEF, France). CH4 and N2 were used as the reactive gases for the experiments.

Glass substrates were used for optical studies. Before being introduced into the deposition chamber, all the substrates were pre-cleaned in ultrasonic bath, in ethanol and were then dried. Two glasses per experiment were prepared. Glasses were positioned back to back on the holder to prevent coating on their back sides. For the profilometer measurements, small lines were drawn using a marker pen on each of the substrates. Coatings on the painted areas were easily removed with ethanol after the experiment to form a sharp thickness difference between the coating and the substrate for thickness measurements.

Table 1. Deposition parameters.

SiCN 1 SiCN 2 SiCN 3 SiCN 4 SiCN 5 SiCN 6 SiCN 7

Base Pressure (Pa)

Working Pressure [Pa]

Ar Gas Flow [cm3/min]

CH4 Gas Flow (cnrVmin]

N2 Gas Flow [cm3/min|

Deposition Temperature [°C]

Sputtering Power [W]

Cathode Voltage [V]

Cathode Current |A|

Bias Voltage |V|

Deposition Time |min|

lxlO"5

0.2

13

5

2

50±10

500

595

0.9±0.1

50

60

lxlO"5

0.2

10

5

5

50±10

500

525

0.9±0.1

50

60

lxlO"5

0.2

8

5

7

50±10

500

576

0.9±0.1

50

60

lxlO"5

0.2

5

5

10

50±10

500

580

0.9±0.1

50

60

lxlO"5

0.2

13

2

5

50±10

500

685

0.9±0.1

50

60

lxlO"5

0.2

8

7

5

50±10

500

568

0.9±0.1

50

60

lxlO"5

0.2

5

10

5

50±10

500

564

0.9±0.1

50

60

A base pressure of lxlO"5 Pa was obtained by a combination of a rotary and turbomolecular pump system before each experiment. Ar (99.999% purity) gas flow was used for generating plasma in the vacuum chamber. Plasma booster was used to increase the plasma density around the substrate holder for bias etching. Bias voltage was gradually increased from 50 V to 250 V and applied for ten minutes during the etching process. After the bias etching plasma booster was turned off, CH4 and N2 were introduced into the chamber as carbon and nitrogen sources in

596

appropriate flow rates. During all coating processes, a total gas flow of 20 cnrVmin was maintained. Both CH4 and N2 flow rates were varied between 1 to 10 cnrVmin. For all the experiments, bias voltage was held constant at 50 V. Depositions were realized without any external heating. The process parameters for SiCN films are listed in Table 1.

2.2. Characterization

Thicknesses of the coatings were measured by a Dektak 150 model profilometer (Veeco, USA). An NKD 7000 model spectrophotometer (Aquila, UK) was used to measure the optical transmittance and reflectance of silicon carbide thin films over the spectral range from 280 to 1000 run at 30° angle of incidence. The optical constants such as refractive indices and the extinction coefficients of the films were calculated using Pro-Optix software. INCA energy dispersive spectroscopy (Oxford, UK) was used for elemental analysis.

3. Results and Discussion

3.1 Profilometer Analysis

Thicknesses of SiCN films were measured by a profilometer. Results showed that all the thicknesses were in the 1350±70 nm range; hence the effects of thicknesses on the optical constants and band gap of SiCN films were negligible. Profilometer measurement results are shown in Table 2.

Table 2. Thicknesses of the deposited films. SiCN SiCN

SiCN 1 SiCN 2 SiCN 3 SiCN 4 SiCN 5 6 7 Film Thickness

|nm| 1380 1420 1390 1370 1370 1320 1280

3.3. Spectrophotometer Analysis

The films deposited under constant 25% CH4 and varied N2 gas flows are called Group 1 and the films deposited under constant 25% N2 and varied CH4 gas flow are called Group 2.

The transmittance and reflectance values of the deposited films were measured by NKD spectrophotometer in the wavelength range of 280-1000 nm. Results of the measurements are shown in Figure 1 and 2.

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Figure 1. Transmittance (T) and reflectance (R) values of Group 1 SiCN thin films under s-polarized waves.

Figure 2. Transmittance (T) and reflectance (R) values of Group 2 SiCN thin films under s-polarized waves.

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As can be seen in Figure 1, transmittance values of Group 1 specimens decreases by the increase in nitrogen gas flow under -550 nm wavelength. However, by increasing the nitrogen gas flow, it can be seen that the transmittance values are increased over -550 nm wavelength. Transmittance values have a tendency of staying constant over -800 nm wavelength. This is a result of the effect of glass which has a maximum transmittance value around 85-90%.

The films which are deposited with 10% and 50% N2 gas flow (SiCN 1 and SiCN 4) are generally more reflective than SiCN 2 and SiCN 3 which are deposited under 25% and 35% N2 gas flow respectively.

As seen in Figure 2, for the depositions realized under constant N2 gas flow, the increase in methane flow rate decreases the transmittance values of the films under -400 nm with the exeption of SiCN2. Below -650 nm wavelength, the increase in carbon concentration in the films results in a dcrease in transmittance values. Above -650 nm, it can be clearly seen that the average trasmittance values can hardly be read because of the interference between waves. Besides this, above -850 nm wavelength, the transmittance values generally increase by the increase in carbon concentration.

Increase in carbon concentration results in a general increase in reflectance values. However, SiCN 5 had the highest value of reflectance throughout the graph.

3.4. Microstructure and Chemical Composition Analyses

Topographic views of the surfaces were observed by scanning electron microscope in order to investigate the microstructural properties of the films. These images are taken for the coatings on M2 steel substrates.

As seen in Figure 3, the deposited films are dense, homogenous and well adherent on the substrate material.

Figure 3. SEM image of SiCN 4 (left) and SiCN 7 (right).

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3.5. Calculation of the Optical Constants and Band Gap

The values of refractive index (n) and extinction coefficient (k) have been calculated by using the theory of reflectivity of light. According to this theory, the reflectance of light from a thin film can be expressed in term of Fresnel's coefficient. The reflectivity [5] of an interface can be given by:

R=[(n-iy+k2]/[(n+l)2 + kz] (1)

The absorption coefficient (a) can be found by the equation [5]:

a = 4nk/X (2)

where, X is the wavelength. Using the transmittance and reflectance data, the extinction coefficient k were evaluated by the spectrophotometer software. The absorption coefficient was calculated by Equation 2 which was later used in Taue method for determining the band gap values.

Empirical expressions for the optical gap can be defined in terms of extrapolation of the extended states. Taue rule is widely used to determine the optical gap in amorphous semiconductors where it is assumed that the square root distributions of conduction and valance band states, and disorder characteristics of amorphous semiconductors relaxes the momentum conservation such that the momentum matrix element is independent of photon energy, hv. Taue, Grigorocivi and Vancu defined an optical gap (£/°uc) that is widely used among experimentalists reporting on the optical properties of thin films [6, 7].

Taue rule was used for the determination of the band gap values of the deposited thin-films. Taue rule [6] can be written as:

>/«(fcr)fcv = CTauc(hv - El"" ) (3)

where, a is the absorption coefficient, hv is the photon energy, CTauc is the Taue constant in this extrapolation, and Ej""' is the corresponding Taue gap [6].

The needed data for Taue method is calculated by the spectrometer software. For the extrapolation method [6], the calculated absorption coefficient in the Taue formula (cdiv)"2 was plotted as function of (hv). With the extrapolation of the linear parts of the graphs, band gaps of coatings were calculated as can be seen in Figure 5. The results of these calculations are shown in Table 4.

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Figure 5. Calculations of band gap values by extrapolation.

4. Conclusions

Homogeneous and well-adherent silicon carbonitride thin films with thicknesses of 1280-1420 nm with variable compositions were successfully deposited on glass substrates. A spectrophotometer with inline software capability was used for determining the transmittance, reflectance and extinction coefficient values of silicon carbide films. Taue rule was used for determining the optical band gap values. Increase in carbon and nitrogen concentration resulted in both increase and decrease in T-R values for different wavelengths. However, it is clear that the increase in these concentrations results in an increase in optical band gap values. The maximum band gap was achieved for the SiC 5 film. The optical gap of silicon carbonitride films are in the range from 2.35 to 3.27 eV.

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Table 4. Band gap values for SiCN thin films.

Gaz Flow [cm /min]

Group 1

Group 2

Specimen

SiCNl

SiCN 2

SiCN 3

SiCN 4

SiCN 5

SiCN 2

SiCN 6

SiCN 7

CH4

5

5

5

5

2

5

7

10

N2

2

5

7

10

5

5

5

5

Band Gap (eV)

3.23

2.63

2.35

2.35

3.27

2.63

2.69

2.58

It can be concluded that carbon and nitrogen concentration in SiCN thin films has significant influence on both optical constants and optical band gaps. The results demonstrated that the transmittance and reflectance values of SiCN films could easily be tailored by modifying Si, C and N concentrations in the coating, for the same film thickness.

5. Acknowledgments

The authors would like to thank Prof. Dr. Fatma Tepehan for the use of profilometer and spectrophotometer, and Prof. Dr. Gültekin Göller for the use of SEM.

References

1. K. B. Sundaram and J. Alizadeh, "Deposition and optical studies of silicon carbide nitride thin films", Thin Solid Films, 370 (2000), 151-154. 2. P. Jedrzejowski et al., "Mechanical and optical properties of hard SiCN coatings prepared by PECVD", Thin Solid Films, 447-448 (2004), 201-207. 3. E. Xie et al., "Preparation and characterization of SiCN films", Optical Materials, 23 (2003), 151-156. 4. A.K.. Costa et al., "Characterization of ultra-hard silicon carbide coatings deposited by RF magnetron sputtering", Thin Solid Films, 377-378 (2000), 243-248. 5. S.A. Khan, M. Zulfequar, and M.Husain, "Optical band gap and optical constants in a-Se8oTe2o-xPbI thin films", Current Applied Physics, 5 (2005), 583-587. 6. A.O. Kodolbas, "Empirical calibration of the optical gap in a-Sii-xCx:H (*<0.20) alloys", Material Science and Engineering, B98 (2003), 161-166. 7. F. Yakuphanoglu, "The effect of FeCh on die optical constants and optical band gap of MBZMA-co-MMA polymer thin films", Physica B, 391 (2007), 136-140.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INFLUENCE OF TIG RE-MELTING AND RE (La2Oj) ADDITION ON MICROSTRUCTURE, HARDNESS AND WEAR OF Ni-WC COMPOSITE

COATING

B.M. Dhakar1, D.K. Dwivedi1*, Satpal Sharma2

'Department of Mechanical & Industrial Engineering, 1.1. T., Roorkee, Roorkee-247667,

Uttranchal, India 2School of Engineering, Gautam Buddha University, Greater Noida, U. P., India.

■"•Corresponding author - E-mail: dkd04fme@iitr.ernet.in, Fax: +91-1332-285665

Keywords: Flame spraying; Abrasive wear, Re-melting, Microstructure.

Abstract

This paper describes the effect of lanthanum oxide (La2Û3) and remelting on microstructure, hardness and abrasive wear behavior of Ni-WC composite powder flame sprayed coatings. It was found that La2Û3 modification and remelting of flame sprayed Ni-base composite coating refined the microstructure and increased microhardness and resistance to abrasive wear. La2C>3 modification and remelting of the coating increased hardness by 1.2 to 1.3 folds. Abrasive wear behavior of unmodified and La2Û3 modified coating in as sprayed and re-melted conditions were studied for different normal loads (5, 10, 15 and 20N) and abrasive mediums of 120 and 600 grit sizes. La2Û3 modified composite coating in re-melted condition showed minimum wear rate irrespective of normal load and abrasive medium. Abrasive wear study results were supported by SEM analysis of worn out surfaces.

1. Introduction

Mechanical failures of components mostly originate from surfaces caused by wear and corrosion. Therefore, to increase the service life of such components, various types of coatings are deposited. For improved surface properties and wear resistance, higher hardness is desirable. It has been reported that the addition of various amounts of Cr-carbide and/or Ti-carbides in Co-based alloys increases the average hardness and wear resistance of composite coatings [1 -5]. An optimum addition of La2Û3 in the TiC/ Ni composite laser clad coatings refines the microstructure, and increases the micro-hardness and wear resistance [6-9], Flame sprayed Fe-Ni coatings with rare earth elements also have been reported to increase the micro-hardness and wear resistance [10]. Wang et al. [11] successfully applied rare earths to modify the surface properties of the thermal sprayed coatings and the laser-alloyed coatings. They studied influence of rare earth elements on the microstructure, wear and corrosion resistance of the modified layers. Refinement in microstructure, increase in micro-hardness and abrasive wear resistance of NiCrBSi flame sprayed coatings with the optimum addition of CeÛ2 (0.8 wt. %) and La2Ch (1.2 wt. %) has also been reported by the author [12]. Literature survey did not reveal any systematic study on the influence of TIG remelting of flame sprayed Ni-WC coating modified by rare earths. Therefore, in the present work, an attempt has been made to study the effect of La2Û3

603

addition and TIG remelting on microstructure, microhardness and wear behavior of Ni-WC composite coatings deposited by flame spraying on mild steel substrates.

2. Experimental procedure

2.1 Materials, methods and characterization of coatings

The nominal composition of substrate and Ni-based surfacing powder are shown in Table I. The commercially available Ni-based composite powder was modified with the addition of 1.2wt.% La2Ü3. The compositions and coating designation are shown in Table II. The flame spraying was carried out using neutral flame of oxy-acetylene gas. The substrate was preheated to 200±10°C. Re-melting of these coatings was carried out using a TIG welding torch operating at DC of 150 A and voltage of 18 V followed by still air room temperature cooling.

Coated samples were cut transversely for microstructural characterization and porosity. Wear test and wear characterization were carried out as described elsewhere [12]. The microstructure of the coating was studied under optical microscope (Leitz, MM6). Vickers hardness of the coating was measured using a load of 100g and average of six readings of the coating has been reported in this study.

Table I Chemical composition (% wt) of substrate and surfacing powder

Material Steel

substrate

Coating powder

B 0.2-0.22

0.4-0.8

Cr -

1.2-1.6

Ni -

Bal.

Si 0.4-0.6

3-3.5

Fe Bal.

_

C -

0.15-0.3

Mn 0.4-0.8

_

WC -

18-20%

Table: II. Designation of various coatings developed.

Composition Coating designation Ni-base composite powder / coating Unmodified powder / coating Ni-base composite powder / coating +1,2%La2Q3 Modified powder / coating

3. Results and discussion

3.1 Microstructure

The microstructures of the unmodified Ni-based composite powder coatings in as sprayed and re-melted condition are shown in Fig. 1 (a, b). The optical micrographs the coatings mainly exhibited Ni solid solution cell (white), fine eutectic (black) and tungsten carbide particles. The porosity was found to be approximately 4%. Re-melting of coating significantly changed the microstructure in respect of grain size. Image analysis of microstructure showed that re-melting of coating reduces average grain size of Ni solid solution cells from 31 to 25 um. The reduction in grain size can be attributed to high cooling rate experienced by coating after TIG melting during solidification.

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The microstructures of 1.2wt.%La2C>3 modified coatings in as sprayed and re-melted conditions are shown in Fig. 2 (a, b). The optical micrographs of as sprayed and re-melted 1.2wt.%La2C>3 modified coatings also exhibited Ni solid solution with fine eutectic mixture in inter-dendritic regions. Image analysis of micrograph of 1.2wt.%La203 modified coating showed average grain size of Ni solid solution to be approx. 25um. TIG remelting of this coating further reduced average grain size of Ni solid solution from approximately 25 to 14 um.

Figure I Optical micrographs of Ni-base Figure 2 Optical micrographs of La2Û3 composite coatings in a) as sprayed and (b) modified Ni-based composite coatings in a) re-melted condition (100X) as sprayed and (b) re-melted condition

(100X)

3.3 Microhardness

The average microhardness (Hvioo) of various as sprayed, modified and re-melted coatings is shown in Table III. It can be seen that remelting of unmodified coatings increased hardness from 393 to 404 HV. La203 modification of as sprayed coating (without remelting) also increased hardness from 393 to 495 HV. Effect of the TIG remelting on the hardness of La2Û3 modified composite coating was found similar to that of unmodified coating. Increase in hardness from 495 HVioo to 534 HVioo after TIG remelting of 1.2wt.% La2Û3 modified coating was observed. Increase in hardness of composite coating due to remelting and La2Û3 modification can be attributed to the refined microstructure and uniform distribution of various elements [12].

3.4 Abrasive wear behavior

Wear rate vs. normal applied load relationships of unmodified and La203 modified coating in the as-sprayed and re-melted conditions are shown in Table IV. It can be seen that, in general, wear rate increased with increase in applied load irrespective of coating and abrasive medium. All coatings showed lower abrasive wear against fine abrasive medium (600 grit size) than coarse abrasive medium (120 grit size). This is due to the fact that, with the increase in load and grit size (120 grit size), the depth of penetration and width of the wear groove increases which results in higher wear rates as compared to low load and small grit size (600 grit size) of the coatings.

Re-melted unmodified and La2Ü3 modified coatings are subjected to lower wear rate than as-sprayed coatings irrespective of normal loads and abrasive medium. The addition of La203 in Ni-based composite coating improves wear resistance. La2Ü3 modification reduced abrasive wear of coatings both in as sprayed and re-melted conditions. Further reduction in the wear rates of the

605

re-melted coatings are mainly due to finer microstructure and higher microhardness of the re-melted coating as compared to as-sprayed coatings. The finer microstructure provides low mean abrasive path which results in lower abrasive wear rate as compared to other coatings. The lower wear rates of re-melted and modified coatings as compared to as sprayed coating are due to higher hardness of these coatings.

Table III: Vickers hardness (VHN) of unmodified and modified powder coating in as-sprayed and re-melted condition

S.N.

1

2

3

4

Coating type and their condition

Ni-base composite powder coating (20%WC), as-sprayed

Ni-base composite powder coating (20%WC), re-melted

Ni-base composite +1.2wt.%La203 modified powder coating, as-sprayed

Ni-base composite +1.2wt.%La2C>3 modified powder coating, re-melted

Average hardness (Hv ioo)

393

404

495

534

4. Scanning Electron Microscopy Analysis

SEM images of worn surfaces under 20 N load and 120 abrasive grit size were analyzed (Fig. 4-5 (a-b)) in order to identify the abrasive wear mechanisms in Ni-based composite and 1.2wt.% La2Û3 modified coatings. The worn surfaces of all the coatings mainly showed the cutting, ploughing and crater as the material removal mechanisms (Fig. 4-5 (a-b)) in these coatings. The weight loss in each coating is determined by the extent of damage caused by these mechanisms. Cutting and ploughing were the main wear mechanisms observed in Ni-based composite as sprayed coating while cutting mechanism was observed in re-melted Ni-based composite coating (Fig. 4 a-b). Ploughing in the Ni-based as sprayed composite coating is due to low hardness of the coating. The abrasive wear marks on wear surfaces of as sprayed coating (tested against high load and coarse abrasive medium) were wider and deeper than that in case of re-melted coatings The wider grooves in unmodified as sprayed coating were due to low hardness as compared to Ni-based re-melted composite coating respectively. Deeper indentation caused by coarse abrasive particles at high load results in more damage on the wear surface than that against low load and fine abrasives. These observations are consistent with wear test results.

Cutting was the main wear mechanism observed in Ni-based composite +1.2wt.% La2Û3 modified powder coating in as sprayed conditions. The cutting grooves are less deep and wider in this coating in re-melted conditions as compared to as sprayed coatings (Fig. 5 a-b). This is attributed to high hardness and fine microstructure of this coating in the as sprayed and re-melted conditions. The Ni-based composite +1.2wt.% La2Ü3 modified powder coating showed the lowest abrasive wear in the re-melted condition as compared to all other coatings.

606

Normal

load, N

5

10

15

20

Un-modified Ni based composite coating

As sprayed

120 grit

807

1544

1887

2685

600 grit

473

854

985

1004

TIG re-melted

120 grit

638

1408

1577

2370

600 grit

401

560

643

825

LaîOî modified Ni based composite coating

As sprayed

120 grit

748

1282

1788

2366

600 grit

384

732

854

967

TIG re-melted

120 grit

586

1134

1535

1892

600 grit

253

568

727

838

Table IV Wear rate (mg/km) vs. applied load (N) relationship for un-modified and La203

modified coating.

Figure 4 SEM micrographs of Ni-based composite coatings against 120 grit size abrasive at 20N load in a) as sprayed and (b) re-melted condition

Figure 5 SEM micrographs of La2Û3 modified Ni-based composite coatings against 120 grit size abrasive at 20N load in a) as sprayed and (b) re-melted condition

5. Conclusions

The following conclusions can be drawn from the present study: 1) The La203 addition refines the microstructure of the coatings as compared to those

without La2Û3 coatings. Remelting further refines the microstructure and reduces average grain size of Ni solid cells from 25 to 14 urn in case of La2Û3 modified coatings.

2) There is an increase in average microhardness by 26% with the addition of La2Ch in Ni-based composite coating in the as sprayed condition. TIG remelting of unmodified and modified coatings further increases the microhardness of these coatings by approximately 32%.

3) The La203 modified coating in re-melted condition showed the highest wear resistance as compared to Ni-based composite coating in the as sprayed condition. The increase in the abrasive wear resistance is due to increase in microhardness of the coating.

4) Cutting, ploughing and craters were the main material removal mechanisms in these coatings.

607

References

1. Nishida, Minora, Effect of TiC-Cr^C2 particles content on abrasive wear resistance of Co-base overlay weld alloy, Journal of the Japan Welding Society, 11 (1993), 156-161.

2. S. Harsha, D.K. Dwivedi, A. Agrawal, Influence of WC addition in Co-Cr-W-Ni-C flame sprayed coatings on microstructure, microhardness and wear behaviour, Surface & Coatings Technology 201 (2007) 5766-5775.

3. A.K. Maiti, N. Mukhopadhyay, R. Raman, Effect of adding WC powder to the feedstock of WC-Co-Cr based HVOF coating and its impact on erosion and abrasion resistance, Surface & Coatings Technology 201 (2007) 7781-7788.

4. A.K.Yadav, N. Arora, and D. K. Dwivedi, On microstructure, hardness and wear behaviour of flame sprayed Co base alloy coating deposited on mild steel, Surface Engineering. 22 (2006) 331-336.

5. S. Hrasha and D.K. Dwivedi, Microstructure, hardness and abrasive wear behavior of flame sprayed Co based alloy coatings, Surface Engineering 23(4) (2007) 1 -6.

6. Li An-min ;Xu Bo-fan; Pan Ying-jun; Effect of IM2O3 on microstructure and property of TiC/Ni-based composite coating; Journal of Iron and Steel Research, 15 (2003) 57-61.

7. Li, Xin; Mo, Yimin; Yu, Xiantao; Liu, Wei; Influence of Laß) on tribological properties of laser clad Ni-base alloy coating for ZL108, Lubrication Engineering, 3(2006), 103-104.

8. K.L. Wang, Q.B. Zhang, M.L. Sun, X.G. Wei, Microstructural characteristics of laser clad coatings with rare earth metal elements, Journal of Materials Processing Technology 139 (2003) 448^152.

9. K.L. Wang, Q.B. Zhang, ML. Sun, X.G. Wei, Y.M. Zhu, Microstructure and corrosion resistance of laser clad coatings with rare earth elements, Corrosion Science 43 (2001)255-267.

10. Zhenyu Zhang, Zhiping Wang, Bunv Liang, Tribological properties of flame sprayed Fe-Ni-RE alloy coatings under reciprocating sliding, Tribology International 39 (2006) 1462-1468.

11. Y. Wang, Q. Zhang, M. Su, Q.P. Zhong, The influence ofCe02 on the corrosion resistance of laser re-melted alloy spray coatings on steel, Scripta Metallurgica et Materialia 32 (6), 1995. 891-894.

12. Satpal Sharma, D K Dwivedi, P K Jain, Effect of IM2O3 addition on the microstructure, hardness and abrasive wear behavior of flame sprayed Ni based coatings, Wear, 267 (2009), 853-859.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EVALUATION OF MECHANICAL PROPERTIES OF Ni-Ti BI-LAYER THIN FILMS

Maryam Mohri1 and Mahmoud Nili-Ahmadabadi1

'School of Metallurgy and Materials Engineering, University of Tehran, Tehran, Iran

Keywords: Bi-layer composite NiTi,Shape memory effect,Super elasticity.Nanostructure.

Abstract

Shape memory thin films are attractive candidates for micro-electro-mechanical-systems because of their large deformation and strong recovery force. In the present study, Ni-Ti thin films have been deposited on glass substrates by dc magnetron sputtering with the source fitted with an alloy target, which was prepared in vacuum arc remelting (VAR). In this study, three types of thin films have been deposited; Ni4sTi5oCus and Niso.7̂ 49.3 thin films were separately deposited on glass substrates and also a composite bi-layer of N150.7T149 3and Ni45Ti5oCujwas separately deposited on glass substrates.

The as-deposited Ni-Ti thin films were crystallized to change from amorphous to nano-structured phase to characterize shape memory and super elastic behaviors. The composition of targets and films were determined by energy dispersive X-ray spectroscopy. The crystallization, surface morphology and structural features were studied using X-ray diffraction (XRD) and atomic force microscope (AFM).Mechanical properties were characterized by nano-indentation and compared to each other.

Introduction

The miniaturization of engineering devices has created interest in new actuation methods that are capable of high power and high frequency responses. Shape memory alloy (SMA) thin films have exhibited one of the highest power densities of any material used in these actuation schemes. Shape Memory Alloys (SMAs) have the unique ability to recover inelastic deformation upon being heated. The effect is due to a diffusion-less first-order solid-solid martensitic transformation which consists of a low temperature martensite phase (monoclinic) and a high temperature austenite phase (cubic) [1-5]. In the form of thin films, SMAs are being developed for a variety of applicationsin micro-electromechanical systems (MEMS). Nickel-titanium (NiTi) SMAs are popular because of their excellent mechanical properties and biocompatibility [6-11]. They have been used to fabricate actuation devices such as micropumps, microvalves and microgrippers [12].

However, they currently requirecomplex thermo-mechanical training in order to be actuated, which becomesmore difficult as devices approach micron scale.Previous studies haveshown that SMA films with compositional gradients have the added feature ofan intrinsic two-way shape memory effect (SME). The bi-layer thin films exhibit a two-way SME with a reducedhysteresis, while the homogeneous films exhibit the classical one-way SME. Most devices that implement

609

SMA films as actuators only permit repeatableactuation behavior by applying a biasing force to homogeneous films.

To further improve the properties of TiNi films, multi-layer, composite or functionally graded TiNi based films can be designed. In order to successfully develop functionally graded TiNi thin films for MEMS applications, it is necessary to characterize, model and control the variations in composition, thermomechanical properties and residual stress in these films [13].

In this work, a NiTi bi layer thin film with austenitic and martensitic layers at room temperature was designed to study its SMA behavior.In this regard, the mechanical properties of the titanium-rich and nickel-rich NiTi monolayers were compared with bi-layer NiTi thin films.

Experimental Details

Two monolayer Ni-Ti and one bi-layer thin films have been deposited on glass substrates by dc magnetron sputtering source fitted with a 80mm diameter alloy target. In order to obtain a variety of film compositions, several discs of alloy targets, which were prepared in vacuum arc remelting (VAR), were used. The deposition pressure was 5xl0"2 mbar argon gas and the thickness of the films was about lum.

The composition of the as-deposited amorphous films was determined by energy dispersive X-ray analysis (EDXA). The results showed thin films with two kinds of compositions Ni45Ti5oCu5and Ni5o,8Ti49.2 (the subscript number index the atomic percentage of alloy elements). The annealing temperature was 500°C for 1 h under vacuum for crystallization of amorphous films.

Nanoindentation was performed using a Hysitron Inc. TriboScope® Nanomechanical Test Instrument and Berkovich diamond indenter was used; loading and unloading were atO. 15 mN/s, giving indentation depths ranging from 10 to 200 nm. Both load and displacement were recorded during the entire loading and unloading cycle.

The post-annealed films were then subjected to nano-indentation analysis. Depth sensing nano-indentation provides the variation in indentation load, P, as a function of indentation penetration depth, h. In the indentation tests, a simple indentation sequence was followed. The indenter was loaded at a specified rate to a pre-determined load and was held constant for 10s at maximum load, followed by unloading. The load-displacement experiments were repeated at five different locations on the surface of films. The surface roughness of the films was also measured using atomic force microscope (AFM: Digital Instruments NanoScope III) prior to nano-indentation measurements. AFM was also used to image the residual indent to determine the extent of pile-up or sink in. In addition, the microstructure of post-annealed films was also investigated by X-ray diffraction to determine the microstructure of each film.

Result and Discussion

XRD Investigation:

Room temperature XRD patterns of Ni45Ti50Cu5 and Ni5o8Ti492 thin film specimens are displayed in Figs, land 2 respectively. As for as-deposited films, their XRD patterns exhibited

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onlya broad peak around 26 = 42.5°, which suggested an amorphous structure. After annealing at 500°C for lh, thin film specimens were completely crystallized.

Based on the XRD analysis, it is revealed that the annealed l^sTisoCusthin film was composed of martensite phase; the annealed Niso.8Ti49 2 thin film was composed of B2 parent phase.

Figure 1. XRD patterns of (a) as deposited and (b) annealedN^TisoCus thin films.

Figure 2. XRD patterns of (a) as deposited and (b) annealedNi50 gTi49 2thin films.

Nanoindentation:

As deposited, the NiTi films are amorphous and the surface topography is relatively uniform. Annealing the films up to 500 °C initiates the crystallization process.

Thin film surface morphology on a large scale scan (50 umx50 urn), after annealing, is dependent on film composition. Fig. 3 shows AFM micrographs of annealed films of varying composition, revealing significant roughness in N^TisoCus (Ra=~20 nm over a scan area of 50 um><50 urn) due to the formation of surface martensite and twinned structures (Fig.3a). The

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surface roughness of an annealed Niso8Ti49 2 film is around 7 nm (see Fig. 3b). As it can be seen from Fig.4, the surface roughness of bi-layer thin film is around 15 nm.

Figure.3. Surface morphology of a) the Ni45Ti5oCu5 films and b) the M50 8T149.2 films.

Figure.4. Surface morphology of the bi layer thin film.

The load versus indentation penetration depth curves for the 500°C annealed is shown in Fig. 5. (All indentation tests were conducted at room temperature). For the Niso 8T149 2films (Fig. 5a), high elastic recovery can be seen in the unloading curve, indicating the predominance of elastic, or pseudo-elastic, deformation. For Ni45Ti5oCu5 films (Fig. 5b), there is much less evidence of such elastic recovery in the unloading curve indicating predominantly plastic deformation. The above measurements show that the Ni4sTisoCus film has the martensite structure at room temperature, whilst the Niso.sTi49.2film is austenitic. Figure 5a in comparison with 5b shows lower indentation depth which is due to the higher strength of austenite in comparison with the martensite phase. In addition, two distinct slopes in the loading curve in Fig. 5a is depicted which refers to the pseudo-elastic behavior of the austenite phase.

For austenite structures, both martensite formation and detwinning can occur, while in the martensitic phase, only detwinning takes place. Martensite formation in the austenitic phase is

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metastable and thus reverses on unloading, leading to recoverable strain, while detwinning results in permanent deformation, unless a subsequent heating step is included [14].

Fig. 6 is a force-displacement curve obtained via nanoindentation. The specimen in this case was an annealed lum NisosTM f̂ilm on a lum Ni45Ti5oCu5 substrate. The indentation depth in this specimen is in the order of austenitic mono-layer. The loading curve in this specimen also shows two different slopes while the second slope is due to a higher load (2000uN) in comparison with lOOOuN in the austenitic mono layer (Fig. 5a). Summarizing the surface roughness and load-unload curves in Fig.6 shows the effect of the bi-layer on the mechanical behavior when it is compared with different mono layers. The curves during loading and lower indentation depth mean that the austenite layer is loaded while the unloading curve shows similar behavior with the martensite phase. Although the indentation depth is lower than the layer thicknesses (1000 nm vs. 150 nm), it seems that the lower layer was also affected during loading.

In fact during the nanoindentationof bi-layer film,the elastic recovery is morecomplex than that of mono layer thin film due to the influenceof interface between the layers and gradation of composition across the thickness. The diffusion process during annealing is anticipated to reduce the level of residual stress even further as an equilibrium state is approached via the compositional gradient.Residual stresses due to the inhomogeneity that exists between the layers may also complicate the indentation response by modifying the resistance of the material to deformation [15].

On the other hand, as a consequence of annealing, the Cu diffused to the austenite phase adjacent to the interface and more Ni diffuses to the martensitic layer which could expand the martensite layer thickness. This phenomenonis due to the fact that, when the loading depth is low, the austenite layer is only deformed. After full loading, the extended martensite phase is affected by the stress field such that after unloading, the martensite prevents full recovery.

Figure.5. Load vs. indentation penetration depth plots for a) the Niso.8̂ 49.2 films and b) NJ45Ti5oCus films.

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Figure.6. Load vs. indentation penetration depth plots forbi-layer thin films.

Hardness, indent depth recovery ratio and dissipation energy of the films have been evaluated by using load-displacement curves in order to reveal the qualitative behavior for the two crystallographic states. The Young's modulus and hardness are calculated from the unloading curve using Oliver and Pharr analysis [16]. It should be noted that the hardness and modulus values obtained through Oliver and Pharr method are overestimated if pile up occurs around the indent. A representative AFM image of the indent along with the indentation profile of the NJ5o.8Ti492 film is shown in Fig. 7. It can be seen that no pile-up has occurred. Thus, the Oliver and Pharr analysis is appropriate for calculation of elastic modulus and hardness in the present case. The values of the hardness and elastic modulus for both the thin films are enumerated in Table I. The hardness value showed that the Niso8Ti492 film shows higher hardness than the Ni45Ti50Cu5 filmwhile the bi-layer values are in between.

TableI.Elastic modulus and hardness for thin films.

Ni 45Ti5„Cu5

Nl50.8T"'49.2

Bi-layer

E(GPa) 17.3 132.1 90.3

Hardness (GPa) 2.2 6.5 4.8

Figure 7. Image and surface profile of the residual indent using AFM for Ni50 8Ti49 2films.

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The indentation induced super elasticity effect can be characterized by the depth recovery ratio from the load versus depth curves by using the following equation [17]:

Depth Recovery Ratio = (hm„~hr) Am„(l)

where,hmax is the penetration depth at the maximum load and hr is the residual depth when the load returns to zero during unloading. The ratio of the (hr) to (hmax) is much lower for the Ni5o.8TL(9.2films as compared to the I^sTisoCusfilms. This indicates that the extent of recovery of the indenter penetration on unloading, which is given by depth recovery ratio (Eq. 1), is higher in the case of the Ni50 sTL»^films as compared to the NJ45TisoCu5 films. This can be attributed to the fact that the microstructure in the Niso sTi49.2films is mostly austenite (B2) phase whereas that for the Ni45Ti5oCu5 films, it consists of martensite phases and it is the B2 phase (austenite) which shows stress induced martensitic transformation (pseudoelastic effect) on indentation loading [18].

The dissipation energy (WD) was evaluated from numerical integration of the area between the loading and unloading curves caused by dislocation movement, martensite transformation and martensite reorientation [19]. Dissipation energy of the film, exhibitingmartensite state at room temperature, was found to be more than the film with austenite state; this could also be due to the shape memory mode of martensite films. In the bi-layer thin films, dissipation energy is between that of martensite and austenite phase.

The indentation induced superelastic energy recovery ratio (nw) was also calculated using the following relation [18]:

f>>mazFdh

vw - Wt ~ I*™* pan

where, We is the reversible work and Wt is the total work done. Super elastic energy recovery ratio at room temperature was found to be 0.74, 0.4 and 0.55 for the Ni50.8^49.2, the Ni45Ti5oCus and bi-layer films, respectively. The recovery value of bi-layer is also between the austenitic and martensitic mono-layer which shows the interference of both layers during loading-unloading.

Conclusions

In the present study, mechanical behavior of two different mono-layers and bi-layer thin films after annealing was studied. The x-ray and nano-indentation results show that:

the Ni45Ti5oCu5 and Niso 8Ti49.2mono-layers thin films show shape memory effect and super elastic behavior at room temperature, respectivelywhile the bi-layer thin film shows mixed behavior;

as a consequence of annealing of bi-layer,Cu at the interface diffuses to the Ni rich layer and vice versawhich lead to a composition gradient. The bi-layer with composition gradient shows a pseudo-elastic behavior during loading and mainly, martensitic behavior during unloading;

the modulus and hardness of graded thin films is less than austenite phase and higher than martensite phase.

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References

1. E. Baldwin, A. Rabiei,"High Transition Temperature Shape Memory Alloys for Micro-actuator Systems",( Paper presented at Materials Research Society Symposium Proceedings, USA, Fall 2003,vol. 785). 2. L. Gian et al.,"Anomalous relationship between hardness and wear properties of a superelastic nickel-titanium alloy ", Applied Physics Letters, 84 (2004), 1076-1078. 3. J. Gill, K. Ho and G. Carman," Three-dimensional thin-film shape memory alloy microactuator with two-way effect", Journal of MicroElectroMechanicalSystems, 11 (2002), 68-77. 4. E. Hornbogen, V. Mertinger and D. Wurtzel," Microstructure and tensile properties of two binary NiTi-alloys",ScriptaMaterialia, 44 (2001), 171-178. 5. W. Tang et al., "New modelling of the B2 phase and its associated martensitic transformation in the Ti-Ni system",ActaMaterialia,47 (1999) 3457-3468. 6. W. Buehler, J. Gilfrich and R. Wiley,"Effect of Low- Temperature Phase Changes on the Mechanical Properties of Alloys near Composition TiNi", Journal of Applied Physics, 34 (1963), 1475- 1477. 7. T. Duerig, A. Pelton and D. Stoeckel, "An overview of nitinol medical applications", Materials Science and Engineering A, 273- 275 (1999), 149-160. 8. K. Otsuka and T. Kakeshita,"Science and Technology of Shape-Memory Alloys: New Developments",Materials Research Society Bulletin, 27 (2002) 91-100. 9. A. Pelton et al.," Nitinol medical devices", Advanced Materials and Processes October, (2005) 63-65. 10. R. Adharapurapu, K. Vecchio," Superelasticity in a New Bioimplant Material: Ni-rich 55NiTi Alloy", Experimental Mechanics, 47 (2007), 365-371. 11. A. Pelton, T. Duerig and D. Stoeckel," A guide to shape memory and superelasticity in Nitinol medical devices",Minimally Invasive Therapy and Allied Technology ,13 (2004), 218-221. 12. Y. Fu et al.," TiNi-based thin films in MEMS applications: a review ", Sensors and Actuators A 112(2004)395-408. 13.Sh. Miyazaki, Y. Q. Fu and W. M. Huang,Thin Film Shape Memory Alloys: Fundamentals and Device Applications(Cambridge University Press 2009). 14. A.J. Muir Wood and T.W. Clyne,"Measurement and modelling of the nanoindentation response of shape memory alloys"*4ctaMaterialia, 54/20 (2006), 5607-5615. 15. Yongqing Fu and Hejun Du, Sam Zhang," Sputtering deposited TiNi films: relationship among processing, stress evolution and phase transformation behaviors", Surface and Coatings Technology, 167( 2003), 120-128. 16.W.C. Oliver and G.M. Pharr,"An improved technique for determining hardness and elastic modulus using load and displacement sensing indentation experiments",Journal of Materials Research, 7 (1992), 1564-1583. 17. L. Qian et al.," Comparison of nano-indentation hardness to microhardness ".Surface and Coatings Technology, 195 (2005), 264-271. 18. W. Ni, Y.T. Cheng, D.S. Grummon," Microscopic superelastic behavior of a nickel-titanium alloy under complex loading conditions", Applied Physics Letters, 82 (2003) 2811. 19. K. Komvopoulos, X.G. Ma," Pseudoelasticity of martensitic titanium-nickel shape-memory films studied by in situ heating nanoindentation and transmission electron microscopy",Applied Physics Letters,87 (2005), 263108.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ANODIC TI0 2 NANOTUBULAR ARRAYS WITH PRE-SYNTHESIZED HYDROXYAPATITE - A PROMISING APPROACH TO ENHANCE THE

BIOCOMPATIBILITY OF TITANIUM

Lu-Ning Wang

Department of Chemical and Materials Engineering, University of Alberta, Edmonton, AB, Canada T6G 2V4

Key words: Titanium dioxide nanotubular arrays, Hydroxyapatite, Implant, Biocompatibility

Abstract

Hydroxyapatite (HA) coating has been widely applied on metallic biomédical implants to enhance their biocompatibility. It has been reported that HA coating can be formed on annealed anodic titanium dioxide nanotubular arrays after immersion in simulated biological fluid (SBF) for about 14 days. In the present study, we apply an alternative immersion method (AIM) to form pre-synthesized HA on TiCh nanotubular arrays. An interdigitation phenomenon of the MG63 human osteosarcoma cells with the TiCh nanotubular arrays with pre-synthesized HA is observed in the in vitro test, indicating excellent biocompatibility and bioactivity for present method. The results demonstrate that the AIM treatment is indeed suitable for the TiÛ2 nanotubular arrays and highly enhance the biocompatibility in comparison with the existing methods, i.e. the annealing of the as-formed TiCh nanotubular arrays.

Introduction

Metallic materials, such as titanium and its alloys, are widely applied in orthopaedic and orthodontic implants because of their excellent mechanical properties, adequate corrosion resistance in order to maintain their function after being subject to an environment which is both hostile and at the same time extremely sensitive [1]. Moreover, titanium and its alloys are bioinert materials, which can be covered by the host organism without being integrated with bone [2]. Thus surface modification on titanium and its alloys was required and has been reported during the last several decades [3], For instance, plasma-sprayed coating of calcium phosphate (CaP), such as hydroxyapatite (HA, Caio(P04)6(OH)2) via plasma spray technique, has been conventional product for clinical applications [3]. However, reports describe the difficulty of controlling the quality, composition, and crystallinity of plasma-sprayed HA coatings [4].

Anodic oxidation is increasingly applied for surface treatment on titanium and its alloys to obtain nanotubular structures at the surface [5], During the last decade, there have been numerous studies on the formation of TiC>2 nanotubular arrays by anodization [6-8]. It was reported that the tubes with diameters ranging between 15-500 nm with up to 1000 microns thickness could be grown under anodic potentials ranging from 5-120 V in F" containing acidic electrolyte, resulting in TiC>2 nanotubular arrays possessing a hollow structure for filling with bioactivating species and providing an interface suitable for anchoring connective tissue [9-12]. It was reported that the TiC>2 nanotubular arrays have many potential biomédical applications, for example, as a bond scale and supporting platform for bone and stem cells, local delivery of antibiotics off-implant at the site of implantation, and the control of hemorrhaging by forming significantly stronger clots with reduced clotting times [13]. Recently, HA coatings on Ti02

nanotubular arrays were developed for biomédical applications [14]. It has been reported that a

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thin layer of HA coating with about 1 urn was formed on amorphous ÜO2 nanotubular arrays after 28 days immersion in simulated biological fluid (SBF). By annealing the arrays at 550 °C, a mixed anatase and rutile phase of TiCh nanotubular arrays may shorten the HA formation period to about 14 days. It has been reported that HA coating was formed on the annealed T1O2 nanotubular arrays by immersion in concentrated SBF (simulated biological fluid with 1.5 times of Ca and P concentration than standard SBF) for 5 days [15]. Another study reported that HA could form on the TiÛ2 nanotubular arrays when it was subjected to immersion in 5 M NaOH at 60 °C for 24 hrs followed by immersion in SBF. However, there was a concern that the as-formed nanotubular structure was dissolved or destroyed during the alkaline treatment and it was difficult to identify the role of the damaged nanotubular structure in the formation of HA [16]. Alternatively, room temperature chemical treatment has also been proposed to induce some calcium phosphates on anodic nanotubular arrays to enhance the biocompatibility of implants. An evaporated-based immersion method was reported to induce amorphous calcium phosphate species on anodic TiCh and Z1O2 nanotubular arrays to obtain HA coating in simulated biological culture. Meanwhile, an alternative immersion method is proposed to induce calcium phosphate coating on anodic Ti02 and Z1O2 nanotubular arrays in order to accelerate the hydroxyapatite formation in simulated biological fluid. Both methods show that the HA coating can be induced on anodic nanotubular arrays very fast within a week [17, 18].

A combination of pre-synthesized HA and anodic T1O2 nanotubular arrays accelerated the formation of HA on TiC"2 nanotubular arrays to enhance the biocompatibility more efficienüy. Consequently, in the present study, we investigated whether anodic IÏO2 nanotubular arrays with pre-synthesized HA showed an enhancement of biocompatibility in cell culture.

Experiments

TiO? nanotubular arrays formation

The sample material was Ti thin foil (99.6% purity, Alfa Aesar, USA) with size of 1.0 cm * 1.0 cm and a thickness of 0.5 mm. Prior to anodization, the samples were ground with sandpaper up to 1200 grit and were ultrasonically rinsed with deionized (01) water and ethanol for 20 min, sequentially, and then dried using compressed air at room temperature. The titanium specimen was rinsed in saturated H2SO4 for 1 min to clean the surface before anodization. Anodization was conducted at room temperature in a two-electrode electrochemical cell with a graphite foil (1 cm x 1cm) used as the counter electrode and 0.5 M H3PO4 + 0.5 wt % NH4F solution as the electrolyte. A direct current power supply (1715A, B&K Precision Corporation, Yorba Linda, CA) was used as the voltage source to drive reactions involved in the anodization process. Potential within 10-20 V was applied over a range of anodization periods. All as-prepared TiC^ nanotubular array films were rinsed using DI water and then dried using compressed air. In order to compare the materials formed using different treatments, some anodically treated samples were annealed at 550 °C for 3 h in air using a furnace with a heating rate of 10 °C/min [19].

Pre-synthesized HA formation

The pre-synthesized HA on TiCh nanotubular arrays was achieved by applying the AIM treatment. The specimens were vertically exposed in the following order: saturated Ca(OH)2, DI water, 0.02 M (NH4)2HPC>4, and DI water respectively. The immersion in each solution lasted 1 min at room temperature and one cycle of immersion was defined as exposure in 3 solutions consequently. The specimens were immersed up to 10 cycles and then rinsed by DI water and dried under compressed air.

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In vitro cell culture evaluation

The MG63 human osteosarcoma cell line was used for in vitro test of biological properties of the HA coating. The cells were cultured in RPMI 1640 medium supplemented with 10% FBS, 100 U/mL penicillin, 100 mg/mL streptomycin, and were incubated at 378 °C in a humidified atmosphere with 5% C02. The samples sterilized in autoclave were inoculated with MG63 cells at a density of 5 x 104 cells/cm2. After 3 days incubation, the samples were gently washed three times with cold 0.1M phosphate buffer, fixed in 2.5% glutaraldehyde in 0.1M phosphate buffer for 2 h at 48 °C, thoroughly washed with 0.1 M phosphate buffer again and dehydrated in serial concentrations of ethanol (30, 50, 70, 90, and 100%). The samples were dried in a critical point drier (HCP-2 HITACHI) and sputtered with a thin gold film and then examined by scanning electron microscope (FESEM LEO 1530).

In order to compare the present method of in vitro cell behaviour with the previous reports, the amorphous and annealed TiC>2 nanotubular arrays at 550 °C were also selected as control samples.

Surface characterizations

Field emission scanning electron microscopy (FE-SEM) was used to determine the surface morphology using a JEOL 6301F SEM with a FE electron source operating at 5 kV. The compositions of the coating were determined with an energy dispersive X-ray spectrometer (EDX). The crystal structure was determined using a thin film X-ray diffractometer (TF-XRD) with Cu K.a radiation. For all me samples after cell culture test, the surface morphologies of all the samples were observed by scanning electron microscopy (LEO-1530, Germany) at 20 kV. Each sample was depicted under a fluorescent microscope (50 and 10 magnification) to obtain the image of the cell distribution.

Results and discussions

Formation of TiO? nanotubular arrays

Fig. 1 shows the top views of Ti02 nanotubular arrays formed at different applied voltage, 20 V (Fig. la), 15 V (Fig. lb), 10 V (Fig. lc) for 150 min. It can be clearly observed that the opening of the nanotube is about 85 nm when the anodized voltage is 20 V. With the decrease in the applied voltage to 10 V, the opening decreases to about 40 nm at 10 V. Fig. Id shows the top view of anodized nanotubular arrays annealed at 550 °C. The morphology shows no significant difference from the as-formed nanotubular arrays. It indicates that the annealing treatment does not destruct the nanotubular arrays structure. Fig. 2a shows the relationship of layer thickness and tube opening with respect to applied voltages. It clearly shows the linear increase in layer thickness and tube opening with respect to the applied voltage, which is in agreement with the previous reports. Fig. 2b summarizes the kinetic curves of the nanotubular layer thickness with respect to the anodization time up to 2.5 hrs. The data show diat the growth of Ti02 nanotubes at all applied voltage increases rapidly in the initial stage, i.e., about 60 min. After the first 60 min, the growth of Ti02 nanotubes gradually decreases and the nanotube length after anodization for 120 min is almost in equilibrium. It is well accepted that the formation of Ti02 nanotubes in F" containing solution relies on three spontaneous processes: field assisted oxidation of titanium to form titanium dioxide, field assisted dissolution of titanium dioxide, and chemical dissolution of the Ti02 by etching with fluoride ions [6-8]. It is believed that the final thickness of nanotube

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array is dominated by the dynamic equilibrium between the oxidation and the dissolution processes. As shown in Fig. 2b, the diffusion of fluoride anions through the short nanotubes to etch the titanium substrate to develop the TiCh nanotubes is relatively facile, as demonstrated in the first 60 min anodization, which can be defined as growth region of nanotubes. However, as the tube becomes longer, the diffusion of F" in the tubes become increasingly difficult, with the consequence that longer time is required for the ions to reach the bottom of TiCh and form TiÛ2 tubes. As a result, the rate of growth of the tubes decreases after 60 min, as shown in Fig. 2b. After about 100 mins, the curves become flatter and the growth region changes to equilibrium region. In order to examine the precalcification and the capability of the HA formation, we choose the layer thickness with 850 nm for the following tests.

Fig.l. FE-SEM images showing top view of TiCh nanotubes after different anodization voltage (a) 20 V, (b) 15 V, (c) 10 V for 2.5 hrs and (d) annealed sample at 550 °C.

Fig. 2. Kinetic curves of TiCh nanotubes under different voltages during 2.5 hrs anodization and (d) Relationship between applied voltage and nanotube thickness and tubular diameter.

Pre-svnthesizing HA on TiO? nanotubular arrays

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Fig. 3 shows the surface morphology and chemical composition of TiCh nanotubular arrays after AIM treatments. Fig. 3a shows the top view of the TiCh nanotubes after AIM treatment for 2 cycles. It can be seen that a few of the openings are covered by tiny particles while the others still keep opening. The insertion illustrates the EDX spectrum in Fig. 3a. It is obvious that a certain amount of Ca and P can be detected. Fig. 3b and the insertion are the top view and the corresponding EDX spectrum of TiÛ2 nanotubular arrays after AIM treatment for 6 cycles. Significant amount of deposition has filled in the nanotubular arrays. By comparison of the EDX spectrum, the relative intensity of Ca and P is higher than that shown in Fig. 3b, which indicates that more Ca and P species deposit on the surface. Fig. 3c shows the top view of TiCh nanotubular arrays after AIM treatment for 10 cycles. The surface is covered by white particles and the nanotubular opening cannot be detected. The insertion clearly shows the Ca and P peaks. Fig. 3d illustrates the deposited Ca amount with respect to the AIM treatment cycles. It is shown that the amount of deposited Ca linearly increases with the cycles. After 10 cycles of AIM treatment, there are about 8 at % of Ca deposited on the nanotubular arrays.

Fig. 3. The surface morphology and EDX spectra of TiÛ2 nanotubular arrays with AIM treatment for (a) 2 cycles, (b) 6 cycles and (c) 10 cycles; (d) The relationship of Ca amount with the AIM treatment cycles

Fig. 4a shows the XRD spectra of Ti foil, as-formed TiÛ2 nanotubular arrays and AIM-treated TiCh nanotubular arrays. It shows the Ti peaks in all three samples. The as-formed Ti02 does not show any peaks assigning to rutile or anatase phase, which indicates amorphous nanotubular arrays after the anodization. After AIM treatment, peaks appear within 26-32°, which is shown in Fig. 4b. By comparison with the standard data (JCPDS 09-0432), it clearly shows that, after AIM treatment, the deposited layer is indeed the crystalline HA.

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Fig. 4. (a) The XRD spectra of Ti, TiC>2 nanotubular arrays and AIM-treated Ti02 nanotubular arrays and (b) The magnification of square in (a) shows the existance of HA.

Evaluation of biological properties

The biocompatibility and bioactivity of the TiC>2 nanotubular arrays with pre-synthesized HA has also been evaluated by in vitro cell culture. The SEM images shown in Fig. 5 illustrate the behavior of the inoculated MG 63 cells on different samples in this study. Fig. 5a shows that the cells attach and grow well on the pre-synthesized HA. By higher magnification, Fig. 5b clearly shows that the cell pseudopods grow along the coating and connect with other cells. A close interaction between cells and coating is discerned, meaning that the HA coating presents a very high affinity to the cells incubated on it. Figs. 5c and 5d show the cells on amorphous TiÛ2 nanotubular arrays. It is clearly shown that the surface coverage by cell decreases and the cells do not connect each other. Figs. 5e and 5f show that the cells attach well on annealed TiÛ2 nanotubular arrays. The cells can attach and spread well on the surface, which is in agreement with previous reports that the annealed sample enhances the biocompatibility. Moreover, by close observation, one can see that cell pseudopods do not grow well along the nanotubular surface. This indicates that cells show less recognition of the surface. Figs. 5g and 5h show the cell on Ti foil. The poor attachment of the cells on Ti foil is clearly seen. In a high magnification, there is no evidence showing the cell interacting with the surface and between each other.

Analyzing adherent cells by fluorescence microscopy after 3 d incubation represents a typical cell spreading of mesenchymal stem cells on anodic T1O2 nanotubular arrays with pre-synthesized HA (Fig. 6a), whereas on annealed samples (Fig. 6b), and amorphous samples (Fig. 6c), a decreased cell spreading after 3 d was found and after 3 d incubation cell growth on Ti foil (Fig. 6d) much less spreading occurred than on the other three samples.

All the experiments of cell culture have demonstrated that the TiÛ2 nanotubular arrays with pre-synthesized HA exhibits an excellent biocompatibility and high bioactivity, which is attributed to the formation of the HA thin layer on the nanotubular arrays. The HA thin layer may provide a closely chemical environment to the cell culture for its spreading and growth.

Conclusions

A series of titanium dioxide nanotubular arrays were fabricated by anodization in F" containing electrolyte from 10 -20 V. The diameter and thickness of the nanotubular arrays varied depending on the applied voltage. The anodization duration did not affect the nanotubular diameter, but only the length of the nanotubular arrays. The tube length was increased in the initial stage of anodization and the rate of growth gradually decreased after about 2.5 h

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anodization. Cell attachment and proliferation test demonstrates that the AIM-treated TiC>2 nanotubular arrays exhibit excellent biocompatibility and high bioactivity, which is attributed to the pre-synthesized HA, which provides a chemical similarity to the biological culture and a well-fed micro environment for cell culturing. These results demonstrate that formation of TiCh nanotubular arrays with pre-synthesized HA via AIM treatment drastically induces the HA deposition and highly enhances the biocompatibility on Ti and Ti-based implant materials.

Fig. 5. SEM micrographs of MG63 cells cultured for 3 d on: Ti02 nanotubular arrays with pre-synthesizing HA (a) and (b), as-formed TiC>2 nanotubular arrays (c) and (d), annealed Ti02 nanotubular arrays (e) and (f) and Ti foil (g) and (h). The pair of two images are magnified at x 1000 and x 5000, respectively.

Fig. 6. Fluorescence microscopy image of MG63 cells on (a) TiCh nanotubular arrays with pre-synthesized HA, (b) as-formed TiC>2 nanotubular arrays, (c) annealed TiOj nanotubular arrays and (d) Ti foil, after incubation for 3 d.

Acknowledgement

This work was supported by Natural Sciences and Engineering Research Council, Canada.

References

1. D. M. Brunette et al., Titanium in Medicine: Materials Science, Surface Science, Engineering, Biological Responses and Medical Applications (Springer-Verlag, Berlin, 2001 ), 171.

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2. H. Tschernitschek, L. Borchers and W. Geurtsen, "Nonalloyed titanium as a bioinert metal — A review,"./. Prosthet. Dent. 96 (1) (2006), 12. 3. K. Duan and R. Wang, "Surface modifications of bone implants through wet chemistry," J. Mater. Chem., 16 (2006), 2309-2321. 4. R. B. Heimann, H. V. Tran and P. Hartmann, "Laser-Raman and Nuclear Magnetic Resonance (NMR) studies on plasma-sprayed hydroxylapatite coatings: Influence of bioinert bond coats on phase composition and résorption kinetics in simulated body fluid," Mat.-wiss. U. Werkstofftech 34 (12) (2003) 1163-1169. 5. D. Gong, C. A. Grimes, O. K. Varghese, W. Hu, R. S. Singh, Z. Chen and E. C. Dickey, "Titanium oxide nanotube arrays prepared by anodic oxidation," J. Mater. Res., 16 (2001) 3331-3334. 6. G. K. Mor, O. K. Varghesea, M. Paulosea, K. Shankar and C. A. Grimes, "A review on highly ordered, vertically oriented TiÛ2 nanotube arrays: Fabrication, material properties, and solar energy applications," Solar Energy Mater. Solar Cells, 90 (14) (2006) 2011-2075. 7. P. Roy, S. Berger and P. Schmuki, "Ti02 nanotubes: synthesis and applications,"y4ngeM>. Chem. Inter. Ed, 50 (2011) 2904-2939. 8. A. Ghicov and P. Schmuki, "Self-ordering electrochemistry: a review on growth and functionality of TiC"2 nanotubes and other self-aligned MOx structures," Chem. Comm. 20 (2009), 2791-2808. 9. S. Oh and S. Jin, "Titanium oxide nanotubes with controlled morphology for enhanced bone growth," Mater Sei Eng C, 26 (2006), 1301-1306. 10. M. Paulose, K. Shankar, S. Yoriya, H. E. Prakasam, O. K. Varghese, G. K. Mor, T. A. Latempa, A. Fitzgerald and C. A. Grimes, "Anodic Growth of Highly Ordered T1O2 Nanotube Arrays to 134 urn in Length," J. Phys. Chem. B, 110 (2006) 16179-16184. 11. S. P. Albu, A. Ghicov, J. M. Macak and P. Schmuki, "250 um long anodic Ti02 nanotubes with hexagonal self-ordering," Phys. Stat. Sol. (RRL), 1 (2007), R65-R67. 12. M. Paulose, H. E. Prakasam, O. K. Varghese, L. Peng, K. C. Popat, G. K. Mor, T. A. Desai, C. A. Grimes, "Ti02 Nanotube Arrays of 1000 urn Length by Anodization of Titanium Foil: Phenol Red Diffusion," J. Phys. Chem. C, 111 (41) (2007) 14992-14997. 13. S. Oh, K. S. Brammer, Y. S. Li, D. Teng, A. J. Engler, S. Chien and S. Jin, "Stem cell fate dictated solely by altered nanotube dimension," Proc. Natl. Acad. Sei., 106 (2009) 2130-2135. 14. J. Park, S. Bauer, K. A. Schlegel, F. W. Neukam, K. von der Mark and P. Schmuki, "Ti02

nanotube surfaces: 15nm—An optimal length scale of surface topography for cell adhesion and differentiation," Small, 5 (2009) 666-671. 15. B. Feng, X. Chu, J. Chen, J. Wang, X. Lu and J. Weng, "Hydroxyapatite coating on titanium surface with titania nanotube layer and its bond strength to substrate," J. Porous Mater., 17 (2010), 453-458. 16. X. Xiao, T. Tian, R. Liu and H. She, "Influence of titania nanotube arrays on biomimetic deposition apatite on titanium by alkali treatment," Mater. Chem. Phys., 106 (15) (2007), 27-32. 17. L. Wang, A. Adams and J. L. Luo, "Enhancement of the capability of hydroxyapatite formation on Zr with anodic Z1O2 nanotubular arrays via an effective dipping pretreatment," J. Biomed. Mater. Res. BApp. Biomater., DOl: 10.1002/jbm.b.31898. 18. A. Kodama, S. Bauer, A. Komatsu, H. Asoh, S. Ono and P. Schmuki, "Bioactivation of titanium surfaces using coatings of TiÛ2 nanotubes rapidly pre-loaded with synthetic hydroxyapatite," Ada Biomater., 5 (6) (2009) 2322-2330. 19. H. Tsuchiya, J. M. Macak, L. Müller, J. Kunze, F. Müller, P. Greil, S. Virtanen, P. Schmuki, "Hydroxyapatite growth on anodic Ti02 nanotubes," J. Biomed. Mater. Res. A, 77A (3) (2006), 534-541.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

PREPARATION AND PROPERTIES OF Cu2ZnSnS4 THIN FILMS BY ELECTRODEPOSITION AND SULFURIZATION

Chao An1, Huimin Lu1, Xi Chen'

'Beihang Univ., School of Materials Sei. & Eng.; 37 Xueyuan Road, Beijing 100191, China

Keywords: Cu2ZnSnS4 thin film; Electrodeposition; Ionic liquid electrolyte; Sulfurization

Abstract

The environment friendly Cu2ZnSnS4 (CZTS) thin film is a promising alternative to semiconductors based on Ga or In as solar absorber material, for its beneficial properties including good optical properties, high absorption coefficient as well as an ideal band gap for photovoltaic applications, not to mention abundant and cheap raw materials. In this paper, Cu2ZnSn4 precursor films were electrochemically deposited on Mo-coated glass substrates at room temperature from the air-and water-stable ionic liquid based on chloride/urea eutectic mixture. Then the precursor films were sulfiiretted in a tube furnace with Argon as the carrier gas at diversified sulfurization temperatures to successfully form CZTS thin films. The films fabricated were characterized by X-ray diffraction (XRD), scanning electron microscope (SEM) and energy dispersive spectroscopy (EDS). The CZTS thin films having good tin pyrite-type phase structure with the preferential grain orientation along (112) were synthesized as revealed in XRD analysis.

Introduction

Cu2ZnSnS4 thin film is one of the most promising solar cell materials for its low cost, promising optical properties (1.4- 1.5 eV band gap energy, high absorption coefficient in the order of 10 cm' '), as well as the abundant and nontoxic constituent elements [1-4]. Theoretical calculations shows that the highest conversion efficiency of CZTS thin film solar cell is 32.3% [5]. In 1997, the Katagiri research team in Nagaoka University in Japan began to conduct research on CZTS thin film solar cells, and the first conversion efficiency was only 0.66% [6], but in 2003, it had reached 5.45% [7], In 2008, the conversion efficiency of cells prepared by radio frequency magnetron sputtering had increased to 6.77% [8]. In 2010, it was reported that, a total area efficiency as high as 7.2% under AM 1.5 illumination and light soaking had been achieved by selenization, tuning the composition of the CZTS nanocrystals and developing a robust film coating method [9]. There remains a big gap between the present conversion efficiency and theoretical expectations. Therefore, it has broad potential and profound significance to perform further research in this regard. In recent years, CZTS thin films have aroused new research mass fervor.

In this article, sulfurization of electrodepositing metal precursors is used to synthesize CZTS thin films. Reline ionic liquid is chosen as the electrolyte. Reline has a wide electrochemical window of over 2.5 V that allows access to the electrodeposition of less noble elements which cannot be plated from aqueous solutions [10]. Because of its low cost and fine stability, Reline has been used in the electrodepositon of many metals.

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In the first section of this paper, the electrochemical behaviour of Cu-Zn-Sn precursor films in Reline was studied. We present the surface morphology and composition of the precursor films. In the second part, the Cu-Zn-Sn precursor films by direct electrodeposition were sulfuretted in adequate sulfur atmosphere in a tube furnace at different sulfurization temperatures; finally, the CZTS thin films were obtained.

Experimental Details

Preparation of Materials

In this experiment, the ionic liquid Reline (99%) was obtained from Shanghai Chengjie Chemical Ltd, China. The anhydrous chloride salt of CuCb, ZnCh and SnCU (98%) was supplied by Alfa Aesar. The mixed solution of 0.2mol/L CuCb, 0.2mol/L ZnCh and 0.4mol/L SnCU in Reline were prepared by adding appropriate amounts of CuCk, ZnCb and SnCU to Reline in Bunsen beaker, and stirring the solutions inside a glove box at 60°C for two hours, until it became homogenous. The solution was maintained at 30°C (±1°C) throughout deposition. The Mo substrates were cut to lcmxlcm, then they were polished, cleaned ultrasonically in acetone for 10 minutes, rinsed in deionized water and dried under vacuum prior to being used in electrodeposition. Sulfur (AR, Beijing Yili Industry of Fine Chemicals Ltd., China) was used without further handling.

Electrodeposition and Sulfurization Experiments

The electrochemical measurements and depositions were carried out in a three electrodes cell, consisting of Mo working electrode, Ft foil counter and Pt wire reference electrode. Cyclic voltammetry was preformed with CHI 604C electrochemical workstation (CH Instrument Co., USA). The deposits were washed in deionized water to remove the residual Reline.

Then the Cu-Zn-Sn precursor films were sulfuretted respectively at 400GC, 500°C and 600°C in a single chamber tube furnace (KTL1600, Nanjing University Instrument Co, China) in sulfur atmosphere to fabricate CZTS films, and the effect of sulfuretted temperature on the properties of films was investigated.

Micrograph and Composition Analyses

The surface appearance of films was investigated by scanning electron microscope (SEM). The composition and crystallographic structure of samples were determined by energy dispersive spectroscopy (EDS) and X-ray diffractometer (XRD).The XRD tests were performed using a Rigaku D/MAX 2500 diffractometer with a Cu Ka source (X=1.5418 Â) radiation. The voltage used was 40 kV with 100 mA current. The SEM and EDS experiments used TESCAN VEGA II -LMU microscope with an accelerating voltage of 20 kV.

Results and Discussion

Electrochemical Analysis

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Generally speaking, when metal chloride dissolves in ionic liquid, it will react with free chloroions, and exists as complex ion in fused salt, shown in the reaction equation (1). So when CuCb, ZnCl2 and SnCU dissolve in ionic liquid Reline, they exist as [CuCl4]

2\ [ZnCl4]2" and

[SnCl6]2". Then, through electrodeposition, they react as shown in the following three equations

(2, 3, 4).

Mcix+pcr->[MCU^F (i) CuCl4

2"+2e-^Cu + 4Cl" (2) ZnCl4

2"+2e"^Zn + 4Cl" (3) SnCl6

2 + 2e" — Sn + 6C1" (4)

In this step, the precursor of Cu, Zn and Sn was synthesized.

Cyclic voltammogram (CV) curves of Reline respectively containing 0.2 M CuCl2, 0.1 M ZnCl2 and 0.1 M SnCl2 on Mo electrode at 30°C with a scan rate of 20mV/s are shown in Figure 1. The ionic liquid keeps electrochemically stable from -0.25V to 0.95V on Pt electrode (inset in Figure 1). In Figure, la, there is a pair of sharp redox waves, reduction peak ciand oxidation peak ai, attributing to copper redox reaction, ci is located at -0.52V, so the constant electrodepositing potential is definite at -0.52 V. This is similar in Figure.lb and Figure.lc. It can be determined that the electrodepositing potential of zinc is -1.0 V and stannum is -0.13 V.

Figure 1. Cyclic voltammogram of the Reline containing (a) 0.2 M CuCl2, (b) 0.1 M ZnCl2 and (c) 0.1 M SnCl2 on Mo at 30°C. Scan rate 20mV/s.

The CV curves of Reline containing 0.2 M CuCl2, 0.1 M ZnCl2 and 0.1 M SnCl2 on Mo electrode with a scan rate of 20 mV/s at 30°C is shown in Figure 2. Scan starts from open circuit potential.

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In Figure 2, there are two pairs of reduction peaks and oxidation peaks, and they are marked out as ai, ci and a2, c2. Contrasted with experiments before, it can be judged that, ai and ci stand for the copper redox reaction, and a2 as well as C2 are associated with the redox reaction of zinc. The redox reaction of stannum is not obvious. According to this situation, the co-deposition potential of Cu-Zn-Sn precursor is selected as -1.0 V.

Figure 2. Cyclic voltammogram of the Reline containing 0.2 M CuCl2, 0.1 M ZnCl2 and 0.1 M SnCl2 on Mo at 30°C. Scan rate 20mV/s.

Figure 3 shows the effect of different scanning rates on the cyclic voltammograms of the Reline containing 0.2 M CuCl2, 0.1 M ZnCl2 and 0.1 M SnCl2 on Mo electrode at 30°C. With the increase in scanning rate, the oxidation peaks move to positive potential obviously, which means that the reduction process is irreversible.

Figure 3. Cyclic voltammograms of the Reline containing 0.2 M CuCl2, 0.1 M ZnCl2 and 0.1 M SnCl2 on Mo at different scanning rates at 30°C.

0.2 M CuCl2, 0.1 M ZnCl2and 0.1 M SnCl2 were deposited on Mo electrode at -1.0V at 30°C. After 30 minutes of deposition, grey metallic luster Cu-Zn-Sn precursor was obtained on Mo electrode. The surface morphologies of Cu-Zn-Sn precursor film in different magnification times are shown in Figure 4. In Figure 4a, the film presents compact globular particles which has size of about l~2um. Figure 4b indicates that the individual crystallite has the size around 0.5um.

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Figure 4. SEM micrographs of the Cu-Zn-Sn precursor electrodeposited on Mo in Reline at -1.0V at 30°C: (a) 3000x; (b) 8000 x.

In terms of EDS analysis in Figure 5 and Table I, the film comprises only Cu, Zn and Sn, without impurity of Reline or chloride. At the same time, it can be known that the atomic percent of Cu, Zn and Sn is 2.37: 1.12: 1, close to the theoretical ratio of 2: 1: 1.

Figure 5. EDS analysis of the area shown in the SEM micrograph of Figure 4a.

Table I. Elemental Composition of the Cu-Zn-Sn precursor

Elements

CuK ZnK SnL

Weight Percent

42.51 19.82 35.67

Atomic Percent

52.78 24.93 22.29

Sulfurization Analasis

Cu-Zn-Sn precursor films were sulfuretted in the tube furnace with Argon flow at a variety of temperatures, and the surface morphologies of obtained CZTS films are presented in Figure 6.

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Figure 6. SEM micrographs of the Cu-Zn-Sn precursor films sulfuretted at various temperatures: (a) 400°C; (b) 500°C; (c) 600 °C.

It can be found that the surface topography of the films is a function of the sulfurization temperature. At 400°C, the film did not have fixed shape. At 500°C, the films exhibited globular shape, and the grain size was about lum. With increasing temperature, the grain grew. At 600 °C, the grain had become agglomerate. Obviously, the most comfortable temperature is 500 °C. EDS analysis indicates the mole ratio of Cu, Zn, Sn and S is similar to theory - 2: 1: 1: 4, as shown in Table II.

Table II. Influence of various temperatures on components of CZTS films

Sulfurization Temperature (T/°C)

400 500 600

Composition / (%)

Cu

25.37 25.29 24.56

Zn

12.36 12.30 12.11

Sn

11.36 11.30 12.11

S

50.91 51.11 51.23

Mole Ratio (Cu: Zn: Sn: S)

2.05: 1:0.92:4.12 2.06: 1:0.92:4.16 2.03: 1: 1: 4.23

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Figure 7 shows the XRD analysis of CZTS films sulfuretted at different temperatures. These figures indicate that, at different temperatures, the films always have characteristic diffraction peaks corresponding to pyrite-type phase of CuîZnSnS.», and they all have the preferential grain orientation along (112).

Figure 7. XRD patterns of CZTS films sulfuretted at different temperatures: (a) 400°C; (b) 500°C; (c) 600°C.

Conclusions

Cu-Zn-Sn precursor was fabricated by electrodeposition in ionic liquid Refine. The atomic percent of Cu, Zn and Sn was close to theory - 2: 1: 1, and the size of the individual crystallite was around 0.5um. Analysis of the effect of different scanning rates on the cyclic voltammograms of 0.2 M CuCl2, 0.1 M ZnCh and 0.1 M SnCb indicated that the reduction process was irreversible.

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CZTS thin films were successfully synthesized by sulfuretting the Cu-Zn-Sn precursor. Sulfurization temperature had significant influence on the surface topography of the films. The grain grew with increasing temperature. At the most comfortable temperature, 500°C, the film presented globular shape, and the grain size was about lum. The mole ratio of Cu, Zn, Sn and S was similar to theory - 2 : 1: 1:4. The films had pyrite-type phase of Cu2ZnSnS4, and had preferential grain orientation along (112).

In order to obtain better results, the components of the films should be controlled precisely in the future.

References

1. H. Katagiri et al., "Characterization of Qi2ZnSnS4 Thin Films Prepared by Vapor Phase Sulfurization," Japanese Journal of Applied Physics, 40 (2001), 500-504.

2. N. Nakayama, K. Itok, "Sprayed Films of Stannite Cu2ZnSnS4," Applied Surface Science, 90 (1996), 171-175.

3. T. Tanaka et al., "Preparation of Cu2ZnSnS4 Thin Films by Hybrid Sputtering," Physics and Chemistry of Solids, 66 (11) (2005), 1978-1981.

4. H. Katagiri et al., "Development of Thin Film Solar Cell Based on Cu2ZnSnS4 Thin Films," Solar Energy Materials and Solar Cells, 65 (2001), 141-148.

5. Q. J. Guo, Hugh W. Hillhouse, and Rakesh Agrawal. "Synthesis of Cu2ZnSnS4 Nanocrystal Ink and Its Use for Solar Cells," Journal of the American Chemical Society, 131 (33) (2009), 11672-11673.

6. H. Katagiri et al., "Preparation and Evaluation of Cu2ZnSnS4 Thin Films by Sulfurization of E B Evaporated Precursors", Solar Energy Materials and Solar Cells, 49 (1997), 407-414.

7. H. Katagiri et al., "Solar Cell Without Environmental Pollution by Using CZTS Thin Film," (Proceedings of the 3rd World Conference on Photovoltaic Energy Conversion, 2003), 2874-2879.

8. H. Katagiri et al., "Development of CZTS-based Thin Film Solar Cell," Thin Solid Films, 517 (2009), 2455-2460.

9. Q. J. Guo et al., "Fabrication of 7.2% Efficient CZTSSe Solar Cells Using CZTS Nanocrystals," Journal of the American Chemical Society, 132 (49) (2010), 17384-17386.

10. AP. Abbott et al., "Selective Extraction of Metals from Mixed Oxide Matrixes Using Choline-Based Ionic Liquids," Inorganic Chemistry, 44 (19) (2005), 6497-6499.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

FORMATION OF CRYSTALLINE AND AMORPHOUS PHASES DURING DEPOSITION OF Ni/Ti,.* THIN FILMS ON SILICON SUBSTRATES -

INTERPRETATION OF EXPERIMENTAL RESULTS USING MOLECULAR DYNAMICS SIMULATIONS

S. Aich1*, B. Geetha Priyadarshini1, M. Gupta1, S. Ghosh1, and M. Chakraborty1'2

'Department of Metallurgical and Materials Engineering, Indian Institute of Technology, Kharagpur 721 302, India 2School of Mechanical Sciences, Indian Institute of Technology, Bhubaneswar 751 013, Orissa,

India

Keywords: Magnetron Sputtering, Crystallinity, Thin film, Texture, Microscopy, Amorphous

Abstract

This research was undertaken to study the crystallization and amorphization of magnetron-sputtered NixTii-x thin films using composition variation and various substrate bias voltages. Variation of deposition efficiency with composition and bias voltage suggests that Ni resists re-sputtering. In the case of Ti thin films, the increase in negative bias voltage caused decrease in crystallinity due to re-sputtering. But, the bias voltage did not affect the crystallinity of Ni thin film significantly. On the other hand, high bias voltage induced crystallinity in otherwise amorphous NixTii_x thin films. In order to explain some of the observed trends, molecular dynamics simulations based on embedded atomic method (EAM) potential for Ni-Ti system were carried out. Simulations could explain the partial amorphization of crystalline films and partial crystallization of amorphous films which occurred due to re-sputtering.

1. Introduction

Magnetron sputtered nickel-based thin film materials are subject matter of scientific interest due to their wide range of applications ranging from contact devices to Li storage materials [1, 2]. Ti-based sputter coatings, however, owing to their wear and corrosion resistance, thermal and chemical stability, and high specific strength have gained much attention as diffusion barrier layers and for orthodontic applications [3]. On the other hand, Ni and Ti thin film deposition has also extended its attention towards the fabrication of Ni-Ti based shape memory alloy thin films for Micro-Electro-Mechanical System (MEMS) based microactuators. The principal requisite for realization of shape memory behavior in Ni-Ti alloy thin films is the near-equiatomic composition of Ni and Ti. Co-deposition from elemental Ni and Ti targets is regarded as the most feasible route to achieve near-equiatomic composition in Ni-Ti thin films since individual target powers can be readily controlled during the sputter deposition process [4]. Also, Miyazaki and Ishida [5] pointed out that the transformation and shape memory characteristics of Ni-Ti thin films are strongly dependent on sputtering conditions (gas pressure, substrate temperature, and substrate bias voltage) apart from other metallurgical factors. Application of substrate-bias voltage causes high energy gain by the ions resulting in re-sputtering from the deposited films leading to formation of vacancies. In particular, the various reports on re-sputtering due to high

" Corresponding author. Tel.: +91-3222-281750; Fax: +91-32222-282280. E-mail Address: saich@.metal.iitkgp.ernet,in (S. Aich)

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ion-bombardment emphasize on decrease in film deposition rate, degradation of film crystallinity, variation in texture, and variation in film composition in alloy films due to the different sputtering yields of the two materials [6, 7].

In recent years, there have been numerous studies that demonstrated the effect of substrate bias voltage on the structural and morphological properties of Ni [8] and Ti thin films [6]. However, the influence of energetic particle bombardment on the microstructure, morphology and texture of co-sputtered Ni-Ti thin films on Si (100) are not well investigated. Therefore, the present study intends to investigate the effect of ion-assisted bias sputtering on the structure and properties of magnetron-sputtered Ni, Ti and NiTi films on Si (100) substrates for films of various composition. The foremost emphasis is laid on the difference in crystallinity, film thickness, and morphology of individually deposited Ni, Ti and co-deposited Ni-Ti films in the presence of substrate bias voltage. Classical molecular dynamics simulations were carried out to understand the structural stability of the thin films under the condition of high negative bias voltage, which is associated with re-sputtering.

2. Experimental

Ni, Ti and Ni-Ti films were deposited using a RF-DC KVT Ltd.magnetron sputtering system which is a three gun sputter down setup. Ni and Ti targets of 99.9% purity were employed. The Ar gas of 99.99% purity was used as the sputtering gas with constant flow rate of 50 seem. The films were deposited on to 3 cm x 3 cm p-type Si (100) substrates. Prior to the depositions, substrates were etched with 2 % HF solution in order to remove native oxide layer. The target to substrate distance was fixed at 125 mm and to maintain uniformity in the film thickness, substrates were rotated at 10 rpm during the entire experiment. The films were deposited at room temperature, with varying substrate bias voltage from 0 to -100 V. The individual target powers for Ni and Ti were varied to obtain near- equiatomic NiTi films.

To estimate the crystalline phases generated and preferred orientations in Ni, Ti, and Ni-Ti films, Grazing Incidence X-ray Diffraction technique (GI-XRD) was employed. The diffraction data were collected using Philips X'Pert diffractometer with Cu-target (X=1.54056Â) and the grazing angle was fixed at 0.5° with scan step size of 0.05° over different 20 angle range. Studies on the microstructure and film growth morphology were done on both planar and cross-sectional films by Carl Zeiss-SUPRA40 Field Emission Scanning Electron Microscope (FE-SEM) with an acceleration voltage of 5 kV. The compositional analysis of NiTi films were carried out using OXFORD instruments INCA Energy Dispersive Spectroscopy (EDS) attached to FE-SEM.

3. Classical Molecular Dynamics Simulations

Classical molecular dynamics (MD) simulations were carried out to understand the effect of vacancy creation as a result of re-sputtering on the stability of the crystal structure. MD simulations have been carried out on LAMMPS (Large Scale Atomistic/Molecular Massively Parallel Simulator) platform in which equations of motion are numerically integrated using velocity-Verlet algorithm. Embedded atom method (EAM) potential developed by Lai et al. [9] was used in the present study.

NiTi crystals with various compositions viz. pure Ni, Ni07Tio3, Nio.6Tio.4, NiosTio2, pure Ti were generated and equilibriated at 300 K. Different percentages of atoms were removed from the crystals, which were subsequently annealed at 100 K and 300 K using NPT ensemble scheme to study the resulting structural changes. The radial distribution function was generated from the phase space at specified intervals of time in order to analyze the structural changes. In

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a similar manner, the effect of re-sputtering on glassy structure was also studied. In this case, the initial glassy structure was generated by melting the crystal and subsequently quenching it.

4. Results and discussion

4.1 Morphological features

Figure 1. FE-SEM images of planar and cross-section (a-b) Ni films deposited at 0 V, (c-d) Ni films deposited at -100 V, (e-f) Ti films deposited at 0 V, (g-h) Ti films deposited at -100 V, (i-j) NiTi films deposited at -50 V (the micronbar in each micrograph is 200 nm).

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Figure 1 shows the FE-SEM micrographs of Ni, Ti, and NiTi films on Si substrate. The presence of fine grains is more apparent only in the case of pure Ni and Ti films. However, crystallinity of the film cannot be ascertained from the micrographs. It can be seen that, with the increase in bias-voltage, the microstructure tends to become more compact and refined. A transition from columnar to fibrous structure was also observed. Above trend can be explained in terms of higher kinetic energy of the deposited atoms leading to higher diffusivity and thus more compaction, in the case of higher bias voltage. Thus, the reason for the appearance of porous structure in unbiased films is the insufficient atomic mobility on the substrate surface during the nucleation and growth of the films. This causes the adatoms to deposit where they land, leading to a fine-grained porous film structure which lies in the zone I (columnar) of Thornton's Structural Zone Model (SZM) [10]. According to revised Structural Zone Model, large substrate bias changes microstructure from zone I to zone T (fine-fibrous columnar) structure [11]. There is an exception to this trend; Ni films deposited at -50 V bias voltage showed more compact structure than the Ni films deposited at higher substrate voltage of -100 V. At this stage, the reason for this exception is unknown. With the increase in the bias voltage, the kinetic energy of the striking adatoms will increase which will lead to high kinetic energy of atoms or temperature at the surface of the deposited films. Enhanced temperature will lead to higher diffusivity as suggested and thus higher adhesion of grains leading to decrease in porosity.

Also, it has been observed from the micrographs that, with the increase in negative bias voltage from 0 to -100 V, the film thickness of pure Ni increased slightly. The same trend was observed for NiTi films. On the other hand, in the case of pure Ti, the film thickness significantly decreased with increase in the negative bias.

An extensive study by Martin et al. [7] revealed that Ti deposition rate decreases with increase in RF bias power from 0 to 600 W at sputtering power of 1000 W. They concluded that the structural modifications and inconsistency in deposition rate depend on the Ar ion energy and the ratio of ion to target atom flux.

The above discussion is found relevant and extended in the present case, stating that the considerable reduction in Ti film thickness is due to increase in the energy of adatoms beyond the threshold energy of sputtering phenomena causing enhanced re-sputtering of adatoms from the film surface.

4.2 Structural Analyses

Figure 2. G1XRD pattern for Ni films deposited under different bias voltage

Figure 3. GIXRD pattern for Ni-Ti films deposited under different bias voltage (a) 0 V, (b) -50 V and (c) -100 V

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The X-ray diffraction patterns of Ni films deposited under various substrate bias voltages are shown in Figure 2. All Ni films are indexed to face-centered cubic (fee) Fm3m phase irrespective of the applied bias voltage. As typically observed in fee metals, a strong <111> fiber texture was observed to be present in Ni films. This texture possibly originates from the (111) plane parallel to the film surface having least surface energy [12]. There is hardly any change in the crystallographic structure of the films with the increase in bias voltage. Zhou et al. [13] observed improvement in the crystallinity of Ni films deposited on Si substrates by DC sputtering when the substrate bias voltage was increased from 0 to -80 V. This was attributed to enhanced surface diffusivity of deposited atoms. However, in the present case, the depositions are carried out at room temperature (30 °C) on Si (100). It has been reported that even at low temperature below -270 °C, formation of amorphous Ni films are not feasible under clean conditions (without any presence of impurities) by sputter deposition process, since it is difficult to seize the Ni adatom movement at that temperature [14].

On the other hand, the structure of NiTi films was more or less amorphous with little extent of crystallinity at higher bias voltage, as shown in Figure 3. In the case of pure Ti (Figure 4), the structure of the films, without any substrate bias, was crystalline and showed strong (100) texture at 29 = 35°. With increase in the bias voltage from 0 V to -100 V, change in the texture and damage of crystallinity was observed. At bias voltage of-50 V, the (100) peak intensity decreased and an enhancedadditional (002) peak was observed at 20 = 38°. Further, increase insubstrate bias voltage to -100 V led to the growth of (101) dominant peak at 20 = 40° and the preferred orientation of Ti grains was changed completely from (002) to (101). Similarly, one of the recent studies by Martins et al. [7] shows how the ion bombardment of the growing film affects the texture evolution in Ni-Ti films grown on SiCVSi substrates. The variation is texture with bias voltage is shown in Figure 5, which shows histogram of texture coefficient (T) defined as follows: (hkt) are (100), (002) or (101) orientations and / is the relative intensity. The

Figure 4. GIXRD pattern for Ti F ig u r e 5- Texture coefficient plots for Ti films deposited under different f l , m s a t various substrate bias voltage bias voltage (a) 0 V, (b) -50 V

observed crystallinity in pure Ni and pure Ti films and amorphous structure in NiTi films can be explained in terms of effect of re-sputtering. Re-sputtering is evident from the variation of film thickness with increase in negative bias voltage. With increase in die bias voltage, the mass flux

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and thus the film thickness should go up, if re-sputtering was absent. Thus the fact that for pure Ni and NiTi films, the thickness does not increase appreciably with increase in bias voltage and for pure Ti it significantly decreases with increase in bias voltage, clearly indicates significant re-sputtering during the deposition. Re-sputtering will randomly eject existing atoms and this phenomenon is expected to have significant effect on the structure. In order to understand the effect of creation of vacancy, classical molecular dynamics simulations were carried out for crystals of different compositions viz., pure Ni, Nio8Tio.2. Nio.7Tio.3, Nio.6Tio.4, and pure Ti. In these crystals, vacancies were created by randomly removing the atoms from the crystal. The vacancy concentration was varied from 5 % to 30%. Simulations were carried out at 300 K and pressure of 1 bar was maintained using NPT (number of atoms, pressure and temperature) ensemble conditions.

The plots of radial distribution functions for different composition and vacancy concentrations are shown in Figure 6. It can be observed that except the case of pure element, the presence of even 5% vacancy led to the formation of amorphous phase. The transition to amorphous structure can be understood in terms of distortion of the crystal structure around a vacancy. This distortion is spontaneous and athermal in nature. Due to the removal of several atoms, the extent of distortion increases leading to formation of amorphous phase. The transformation is athermal because it takes place within a time span of few picoseconds. To confirm the athermal nature of the transformation, some of the simulations were repeated at much lower temperatures.

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Figure 6: Radial Distribution Function as a result of annealing at 300 K of NiTi system with different concentrations of vacancies

Even at temperatures as low as 100 K, formation of glassy phase was observed, as shown in Figure 7. According to the simulation results, amorphization of pure elements Ti and Ni is difficult, as larger concentration of vacancies is required. The predicted trend, i.e., stability of crystal structure in pure Ni and Ti films, matches well with the experimental observations.

An otherwise amorphous structure formed as a result of re-sputtering may also transform into crystalline phase, as might have possibly happened in the case of NiTi film at a bias voltage of-100 V. In order to understand this transition, MD simulations were carried out. After creating an initial NiTi glassy structure, vacancies were created by random removal of atoms followed by subsequent annealing. As shown in Figure 8, a small peak appears indicating ordering. This transformation mechanism cannot be captured properly by MD simulation as the precipitation of crystal, being a thermally activated process, is relatively a very large time scale process.

Figure 7. Radial Distribution Function Figure 8: Radial Distribution Function as a result of annealing at 100 K. of NiTi as a result of annealing at 300 K of NiTi system with 10% vacancy glass system with 20% vacancy

With increase in negative bias, the micro-strain in crystal (along the c-axis) was found to vary. At 0 V bias, the Ti films exhibited micro-strain which was tensile in nature. With increase in bias voltage to -50 V, the micro-strain became negative (compressive). The microstrain again became positive (tensile) with further increase in substrate bias voltage to -100 V.

5. Conclusions

The composition and bias voltage were found to have significant influence on the formation of crystalline/amorphous phases as well as the texture of the NixTii_x films. The formation of crystalline structure in pure Ti and Ni films was explained using classical molecular dynamics (MD) simulations. At higher negative bias voltage, significant resputtering occurred. However, resputtering did not affect the stability of theory stall ine structure in pure Ti and Ni films. The formation of crystalline structure in pure Ti and Ni films was explained using classical MD simulations. In the case of NiTi thin films, although it was amorphous without using any bias voltage, little crystallinity was observed when the negative bias was increased. This trend was also explained using MD simulations.

Acknowledgements

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The authors are grateful to the Indian Institute of Technology, Kharagpur for the financial support for this research.

References

[I] P.L. Tarn, and L. Nyborg, "Sputter Deposition and XPS Analysis of Nickel Suicide Thin Films", Surf. & Coat. Technology, 203 (2009), 2886.

[2] L. Ai et al., "Influence of substrate temperature on electrical and optical properties of p-type semitransparent conductive nickel oxide thin films deposited by radio frequency sputtering", Appli. Surf. Sei., 254 (2008), 2401.

[3] V. Chawla, R. Jayaganthan, A.K. Chawla, and R. Chandra, Morphological study of magnetron sputteredTi thin films on silicon substrate, Mater. Chem. Phys., I l l (2008), 414.

[4] S. Inoue, N. Sawada, and T. Namazu, "Effect of Zr content on mechanical properties of Ti-Ni-Zr shape memory alloy films prepared by dc magnetron sputtering", Vacuum, 83 (2009), 664.

[5] S. Miyazaki, and A. Ishida, "Martensitic transformation and shape memory behavior in sputter-deposited TiNi-base thin films", Mater. Sei. Engg. A, 106 (1999), 273-275.

[6] M. Naoe, S. Ono, and T. Hirata, "Crystal orientation in titanium thin films deposited by the sputtering method without plasma damage", Mater. Sei. Engg., A134 (1991), 1292.

[7] R.M.S. Martins et al., "Growth of sputter-deposited Ni-Ti films: Effect of a Si02 buffer layer", Appl. Phys. A, 84 (2006), 285.

[8] R. Mitra et al., "Effect of process variables on the structure, residual stress, and hardness of sputtered nanocrystalline nickel films", J. Mater. Res., 16 (2001), 1010.

[9] W. S. Lai, B. and X. Liu, "Lattice stability of some Ni-Ti alloy phases versus their chemical composition and disordering", J. Phys.:Condens. Matter., 12 (2000), L53-L60.

[10] J. A. Thornton, "Influence of apparatus geometry and deposition conditions on the structure and topography of thick sputtered coatings", J. Vac. Sei. Technol, 11(1974), 666.

[II] R. Messier, A. P. Giri, and R. A. Roy, "Revised structure zone model for thin film physical structure", J. Vac. Sei. Technol., A 2 (1984), 500.

[12] B. Geetha Priyadarshini, S. Aich, and M. Chakraborty, "Structural and morphological investigations on DC-magnetron-sputtered Ni films deposited on Si(100)", J. Mat. Sei., 46 (2011), 2860-2873.

[13] R. Zhou et al., "Characterization of the interfacial reaction between sputter-deposited Ni film and Si substrate", Appl. Phys. A, 80 (2005), 179.

[14] J. G. Wright, "Amorphous Transition Metal Films", IEEE Transactions on Magnetics, 12 (1976), 95.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DOPING AND CO-DOPING OF BANDGAP-ENGINEERED ZnO FILMS FOR SOLAR DRIVEN HYDROGEN PRODUCTION

Sudhakar Shet,1,2Nuggehalli Ravindra,2 YanfaYan,1 andMowafakAl-Jassim1

1 National Renewable Energy Laboratory, Golden, CO 80401 USA 2 New Jersey Institute of Technology, Newark, NJ 07102 USA

Keywords: co-doping, photoelectrochemical, thin film

Abstract

Co-doped ZnO:(Al,N) and ZnO:(Ga,N) films were deposited by co-sputtering using radio-frequency magnetron sputtering on F-doped tin-oxide-coated glass. We found that the ZnO:(Al,N) and ZnO:(Ga,N) films exhibited greatly enhanced crystallinity compared to ZnO:N films doped by pure N and deposited under similar conditions. Furthermore, the ZnO:(Al,N) and ZnO:(Ga,N) films showed much higher N-incorporation than ZnO:N films deposited with pure N doping. As a result, the ZnO:(Ga,N) films showed significantly higher photocurrents than ZnO:N doped only by N. The ZnO:(Cu,Ga) films were synthesized by RF magnetron sputtering in O2 gas ambient at room temperature and then annealed at 500°C in air for 2 hours. We found that the carrier concentration tuning does not significantly change the bandgap and crystallinity of the ZnO:Cu films. However, it can optimize the carrier concentration and thus dramatically enhance PEC response for the bandgap-reducedp-type ZnO thin films.

Introduction

Transition-metal oxides are potential candidates for photoelectrochemical (PEC) H2 production from water [1,2]. However, to date, only T1O2 has received extensive attention. ZnO has similar bandgap and band-edge positions compared to TiÛ2, but ZnO has a direct bandgap and higher electron mobility than TiÛ2. Thus, the PEC property of ZnO also needs to be studied. Like TiÛ2, the bandgap of ZnO (3.3 eV) is too large to effectively use visible light. Therefore, it is critical to reduce the bandgap of ZnO. To date, impurity incorporation has been the main method to reduce the bandgap of Ti02. It has been reported that N-, C-, and S-doping can successfully narrow the bandgap of Ti02 and push the photoresponse into the long-wavelength region. Significant amounts of N can be incorporated into ZnO and W03 only at low temperatures. However, films deposited at low temperature usually exhibit very poor crystallinity, which is extremely detrimental to PEC performance. This dilemma hinders the PEC performance of N-incorporated ZnO and WO3 films. A possible cause for the inferior crystallinity may be uncompensated charged N atoms. This problem could be overcome by charge-compensated donor-acceptor doping, such as co-doping ZnO with (Ga and N), (AI and N) and (Cu and Ga). The effect of passive co-doping of (Ga and N) (AI and N) and (Cu and Ga) codoped ZnO films on PEC performance has not been investigated [3-5].

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In this paper, we report on the synthesis of ZnO:(Ga,N) and ZnO:(Al,N) thin films by reactive radio-frequency (RF) magnetron sputtering in mixed N2 and O2 ambient with low O2 mass flow rate. We found that charge-compensated donor-acceptor co-doping exhibits enhanced crystallinity and incorporates much higher N concentration in ZnO:(Al,N) and ZnO:(Ga,N) thin films as compared to ZnO doped solely with N. As a result, ZnO:(Ga,N) and ZnO:(Al,N) thin films presented improved PEC response, compared to ZnO:N(l) and ZnO:N(2) doped by N alone at both high and low temperatures. We have also synthesized bandgap-reduced /»-type ZnO thin films with controlled carrier concentration through Cu and Ga co-doping. The incorporation of Cu is responsible for bandgap reduction, and the incorporation of Ga is responsible for carrier concentration tuning. We found that the carrier concentration tuning does not significantly change the bandgap and crystallinity. However, it can optimize the carrier concentration and thus dramatically enhance PEC response for the bandgap-reduced p-type ZnO thin films.

Experimental Details

Three sets of experiments were carried out. In the first set of experiments, Ga and N doped ZnO:(Ga,N) thin films were deposited. Ga2Û3 powders (0.007 g) were dispersed uniformly on the Zn metal target of 3-inch diameter as the Ga source. For a comparison, N doped ZnO:N(l) film was deposited. The base pressure was below 5*I0"6 torr, and working pressure was 2* 10"2 torr. The sputtering ambient was mixed N2 and 0 2 with an oxygen gas ratio 02/(N2+02) = 2.5%. Two sets of samples, ZnO:(Ga,N) were grown with RF powers of 100 W at RT and annealed at 500°C in air for 2 hours and ZnO:N(l) film was deposited at substrate temperature of 500°C. All samples were controlled to have a similar film thickness of about 500 nm as measured by stylus profilometry. In the second set of experiments, Al and N doped ZnO:(Al,N) thin films were deposited. For a comparison, N doped ZnO:N(2) film was deposited using ZnO target. Sputtering was conducted with RF powers of 300 W at substrate temperature of 100°C. F-doped Sn02

(FTO, 20-23 fi/D)-coated transparent glasses were used as substrates. The substrate was rotated at 30 rpm for uniform deposition of the film. Prior to sputtering, a pre-sputtering process was performed for 30 min to eliminate any contaminants from the target. For the third set of samples, ZnO:(Cu,Ga) films were deposited using a reactive RF magnetron sputtering system at room temperature and annealed at 500°C in air for 2 h. Conducting transparent fluorine-doped tin oxide (FTO) (20-23 fi/D)-coated glass was used as the substrate to allow PEC measurements. A Cu source in the form of Cu chips (area: 2><5 mm2) were dispersed uniformly on the 3-inch-diamter Zn target. Based on our previous results, ten Cu chips were used to produce bandgap-reduced p-type ZnO. Gallium oxide (Ga2Û3) powders were dispersed uniformly on the target as a Ga source. Two sets of ZnO:(Cu,Ga) films were deposited: a low-Ga regime with Ga203 powders of 0.001 and 0.002 grams, and a high-Ga regime with Ga203 powders of 0.01 and 0.03 grams. We refer to these samples as ZnO:(Cu,Ga)0.001, ZnO:(Cu,Ga)0.002, ZnO:(Cu,Ga)0.01, and ZnO:(Cu,Ga)0.03. The samples in the first set are 0.5 urn thick, and the samples in the second set are 1 urn thick.

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The structure of synthesized films was characterized by X-ray diffraction (XGEN-4000, Scintag Inc.) operated with a Cu Ka radiation source at 45 kV and 37 mA. The N concentration in the thin films was evaluated by X-ray photoelectron spectroscopy (XPS). The ultraviolet-visible (UV-VIS) absorption spectra of the samples were measured by an n&k analyzer 1280 (n&k Technology, Inc.) to investigate the optical properties of deposited thin films. Photoelectrochemical measurements were performed in a three-electrode cell with a flat quartz-glass window to facilitate illumination to the photoelectrode surface. The sputter-deposited films were used as the working electrodes with an active surface area of about 0.25 cm2 [6-11]. Pt mesh and an Ag/AgCl electrode were used as counter and reference electrodes, respectively. A 0.5-M Na2S04 aqueous solution with a pH of 6.8 was used as the electrolyte for the PEC measurements [3-11].

Results and Discussion

Figure 1 shows the XRD curves of the samples of the first set of experiments. The co-doped ZnO:(Ga,N) films exhibited significantly enhanced crystallinity, yet with the incorporation of N. From the XRD curve shown in Fig. 1, the crystallite size of ZnO:(Ga,N) films is estimated to be around 40 nm, which is better than the ZnO:N(l) films grown at 500°C.

Figure 1. XRD curves for a ZnO:N(l) and ZnO:(Ga,N) film, respectively.

The crystallite size of ZnO:N is estimated to be around 29 nm. The crystalline quality was characterized from the full width at half maximum (FWHM) of the XRD curve instead of the peak intensity, because the FWHM is an intrinsic property. Such a significantly enhanced crystallinity could be attributed to the Ga source used (i.e., Ga2Û3). Because the ionic radius of substitutional Ga3+ is close to that of Zn2+, the size mismatch is very small. Furthermore, the Ga203 source provides more oxygen during the sputtering process, leading to decreased oxygen vacancies and enhanced crystallinity. The nitrogen concentration in the ZnO:(Ga,N) films was estimated to be about 2 at%, whereas, no nitrogen is present in the ZnO:N(l) films deposited at 500°C confirmed by XPS measurements. ZnO films with high crystallinity can be synthesized at high temperatures; but at these substrate temperatures, it is very difficult to incorporate a significant amount of nitrogen into the films.

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From the optical absorption coefficients for the samples used in this experiment, the bandgap of the ZnO:N(l) films deposited at 500°C is 3.27 eV, which corresponds well to the bandgap of the pure ZnO. The co-doped ZnO:(Ga,N) films exhibited a bandgap of 3.19 eV, which is slightly lower than that of the ZnO:N(l) films deposited at 500°C. This small reduction in bandgap is due to N-induced upshifting of the valance band. It is shown theoretically that the incorporated N would generate an impurity band above the valance band. Figure 2 shows the photocurrent-voltage curves of ZnO:N(l) and ZnO:(Ga,N) films under illumination with the UV/IR filter.

Figure 2. Photocurrent-voltage curves of a ZnO:N(l) and ZnO:(Ga,N) film, respectively, under the illumination with a UV/IR filter.

It shows clearly that the ZnO:(Ga,N) films exhibited significantly increased photocurrents, compared to the ZnO:N(l) films. At the potential of 1.2 V, the photocurrents were 484 and 38 jiAcm"2 for the co-doped ZnO:(Ga,N) and ZnO:N(l) films, respectively.

Figure 3 shows the XRD curves of the samples of the second set of experiments. The dotted lines in the XRD plot indicate substrate peaks. It is clearly shown that Al and N co-doped ZnO:(Al,N) films exhibited significantly enhanced crystallinity compared to N doped ZnO:N(2) films. Such significantly enhanced crystallinity is attributed to the charge-compensated donor-acceptor codoping. Applying the Debye-Scherrer equation to XRD data, crystallite sizes were estimated to be 24 and 39 nm for the ZnO:N(2) and ZnO:(Al,N) films, respectively. The N concentrations of ZnO:N(l) and ZnO:(Al,N) films were about 1 and 5 at.%, respectively, as determined by XPS.

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Figure 3. XRD curves for ZnO:N(2) and ZnO:(Al,N) films, respectively.

From the optical absorption coefficient of the ZnO:N(2) and ZnO:(Al,N) films, it is seen that the ZnO:(Al,N) films showed optical absorption in a much longer-wavelength region as compared to ZnO:N films, indicating that a significant amount of N is incorporated in the films. The direct optical bandgaps measured for ZnO:(Al,N) films drastically reduced from 3.22 to 2.02 eV, respectively. This significant bandgap reduction is due to enhanced N concentration incorporated in the films.

Figures 4 shows the photocurrent-voltage curves of the ZnO:N(2) and ZnO:(Al,N) films, under illumination with the UV/IR filter. It clearly shows that the ZnO:(Al,N) films exhibited enhanced photocurrents, compared to the ZnO:N(2) films. At the potential of 1.2 V, the photocurrents were 13.6 and 54.3 uAcm-2 for the ZnO:N(2) and ZnO:(Al,N) films, respectively. It indicates that a very high recombination rate of the photogenerated electrons and holes is present in the ZnO:N(2) films, due to its inferior crystallinity, wide bandgap and uncompensated charges. On the other hand, the co-doped ZnO:(Ga,N) and ZnO:(Al,N) films exhibited remarkably increased crystallinity, reduced bandgap and good charge compensation, which lead to enhanced photocurrent than the ZnO:N films.

Figure 4. Photocurrent-voltage curves of ZnO:N(2) and ZnO:(Al,N) films respectively, under illumination with a UV/IR filter.

The results demonstrate clearly that significantly reduced bandgap and enhanced photocurrents can be obtained from charge compensated co-doping approach.

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For Cu and Ga doped ZnO films, we first show the results obtained from the first set of samples. Figure 5 shows XRD curves for the first set of samples.

Figure 5. XRD curves of ZnO:(Cu,Ga)0.001, ZnO:(Cu,Ga)0.002, ZnO:Cu, and ZnO samples.

For comparison, XRD curves of a pure ZnO sample and a ZnO:Cu sample with similar thickness are also shown. The location of the (111) peak of Ga2Û3 is indicated by the red line. It is seen that the incorporation of Cu leads to decreased crystallinity as compared to pure ZnO films.

The incorporation of additional Ga showed similar crystallinity as compared to ZnO:Cu films. The XRD curves showed no Cu oxides and Ga oxides, thus indicating no phase separation. However, the incorporation of Cu and Ga changed the preferred orientation of the films. For example, with the dopants, the (002) peak was suppressed and the (100) and (101) peaks were enhanced. AFM images (not shown here) also confirmed that the ZnO:Cu and ZnO:(Cu,Ga) films have similar grain sizes. It indicates that the incorporation of a low concentration of Ga did not significantly change the morphology and average crystallite sizes.

Figure 6 shows the measured photocurrent-voltage curves under illumination with the UV/IR filter and dark currents for ZnO:Cu, ZnO:(Cu,Ga)0.001, and ZnO:(Cu,Ga)0.002. It clearly shows the effect of co-doping of Ga in improving the PEC performance of ZnO:Cu films. Both ZnO:(Cu,Ga)0.001 and ZnO:(Cu,Ga)0.002 films showed p-type conductivities, indicating that the role of Ga in these films is to reduce the hole concentration that is generated by Cu incorporation.

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Figure 6. Photocurrent-voltage curves under (arrowed) continuous illumination, (black curve) dark condition, with a UV/IR filter measured from ZnO:Cu, ZnO:(Cu,Ga)0.001,

and ZnO:(Cu,Ga)0.002 films.

The reduction in hole concentration increases the depletion width so that more photon-generated electron-hole pairs can be collected. Furthermore, the incorporation of Ga will neutralize charged Cu ions, so that recombination will also be suppressed. Thus, ZnO:(Cu,Ga)0.001 and ZnO:(Cu,Ga)0.002 films showed significantly increased photocurrents as compared to ZnO:Cu films. The photocurrent at a potential of -0.5 V for the ZnO:(Cu,Ga)0.001 sample improved by almost one order of magnitude. The ZnO:(Cu,Ga)0.002 sample showed lower photocurrents than the ZnO:(Cu,Ga)0.001 sample, indicating that an optimum carrier concentration exists The incorporation of more than a certain level of Ga may even cause p-type ZnO:Cu film to become w-type. This is exactly the case for our second set of samples. Figure 7 shows XRD curves for the second set of samples. For comparison, XRD curves of a pure ZnO sample and a ZnO:Cu sample with similar thickness are also shown. All the peaks have higher intensities than that of the first set of samples due to larger thickness. Similar to the first set of samples, the incorporation of Cu changes the preferred orientation as compared to ZnO. The incorporation of additional Ga did not significantly change the crystallinity as compared to ZnO:Cu films. The XRD curves showed no Cu oxides or Ga oxides, thus indicating no phase separation. However, the incorporation of Cu and Ga changed the preferred orientation of the films.

Figure 7. XRD curves of ZnO:(Cu,Ga)0.01, ZnO:(Cu,Ga)0.03, ZnO:Cu, and ZnO samples

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As in the first set of samples, blue shifts are also observed for the second set of samples; however, the shift is much more significant. This indicates that more unoccupied states are now filled by the electrons from Ga, as more Ga2C>3 powders were dispersed on the target. The ZnO:(Cu,Ga)0.03 sample exhibits a larger blue shift than the ZnO:(Cu,Ga)0.01 sample, due to its higher Ga concentration. Figures 8 shows the measured photocurrent-voltage curves under illumination with the UV/IR filter and dark currents for ZnO:(Cu,Ga)0.03 films. We see that ZnO:(Cu,Ga)0.03 films are no longer p-type, but are «-type. ZnO:(Cu,Ga)0.01 films also showed n-type behavior (not shown here). Thus, the Ga concentration is so high in these samples that holes are completely compensated and the films become n-type. Both ZnO:(Cu,Ga)0.01 and ZnO:(Cu,Ga)0.03 films showed very high dark currents indicating poor quality, which may be the result of incorporation of too much Ga.

Figure 8. Photocurrent-voltage curves under (red curve) continuous illumination, (black curve) dark condition, with aUV/IR filter measured from ZnO:(Cu,Ga)0.03 films.

Because the photocurrent-voltage curves indicate that both ZnO:(Cu,Ga)0.01 and ZnO:(Cu,Ga)0.03 are «-type, the bottom of the conduction bands of these materials may be filled with electrons, which partially contribute to the large blue shift.

Conclusions

Nitrogen doped ZnO films, ZnO:N(l) and ZnO:N(2) were deposited at high and low substrate temperatures of 500 °C and 100 °C respectively. Co-doped ZnO:(Ga,N) and ZnO:(Al,N) films were deposited by co-sputtering on FTO at room temperature and 100 °C, followed by post-annealing at 500 °C. We found that, the ZnO:(Ga,N) and ZnO:(Al,N) films showed much higher N-incorporation than ZnO:N(l) and ZnO:N(2) films deposited at 500 °C and 100 °C with pure N doping. Furthermore, the ZnO:(Ga,N) and ZnO:(Al,N) films exhibited greatly enhanced crystallinity compared to ZnO:N(l) and ZnO:N(2) films doped by pure N. As a result, the ZnO:(Ga,N) and ZnO:(Al,N) films showed significantly higher photocurrents than ZnO:N(l) and ZnO:N(2) doped by N alone at both low and high temperatures. For the ZnO:(Cu,Ga) films, carrier concentration tuning does not significantly change the bandgap and crystallinity of the ZnO:Cu films. However, it can optimize the carrier concentration and thus dramatically enhance PEC response for the bandgap-reducedp-type ZnO thin films.

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Acknowledgements

This work was supported by the U.S. Department of Energy.

References

[I] M. Gratzel Nature 414, (2001), p. 338. [2] T. Bak, J. Nowotny, M. Rekas, and C.C. Sorrell, International J. of Hydrogen

Energy 27, (2002), p. 991, [3] S. Shet, K. -S . Ahn, T. Deutsch, H. Wang, N. Ravindra, Y. Yan, J. Turner, M. Al-

Jassim, J. Mater. Research 25, Doi: 10.1557/JMR.2010.0017, (2010), p. 69. [4] K.-S. Ahn, Y. Yan, S. Shet, T. Deutsch, J. Turner, and M. Al-Jassim, Appl. Phys.Lett.

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Mater. Science DOI 10.1007/sl 0853-010-4561-x, (2010). [6] S. Shet, K. -S . Ahn, Y. Yan, T. Deutsch, K. M. Chrusrowski, J. Turner, M. Al-Jassim,

and N. Ravindra, J. Appl. Phys. 103, (2008) p. 073504. [7] K. -S . Ahn, S. Shet, T. Deutsch, C. S. Jiang, Y. Yan, M. Al-Jassim, and J. Turner, J.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MAGNETIC FIELD ASSISTED HETEROGENEOUS DEVICE ASSEMBLY

Vijay Kasisomayajula, Michael Booty1,2, Anthony T. Fiory1, Nuggehalli M. Ravindra1

'Dept. of Physics, New Jersey Institute of Technology, University Heights, Newark, NJ 07102, USA

department of Applied Mathematics, New Jersey Institute of Technology, University Heights, Newark, NJ 07102, USA

Magnetic Fields, Device Assembly, Deterministic

Abstract

Heterogeneous device assembly using magnetic fields is analyzed. Small cylindrical solenoids are used as sources for a controlled magnetic field. This paper investigates the theoretical and experimental lower limit using technology in extant on the dimension of the solenoids and the devices that can be assembled using magnetic field assisted assembly.

Introduction

Background: In recent years, magnetic field assisted assembly of sub millimeter to micron size devices has attracted significant attention due to the heterogeneous nature of the devices and the substrates. The motivation behind this is the ever shrinking world of electronics and the demand for versatility of the devices. The need for a single hand held device to perform a multitude of tasks is sought not only in the telecommunications industry (cellular phone and related mobile devices) but also in the medical field, environmental protection, defense sector to name a few. These devices integrate many components into a very small form factor which demands an extremely efficient assembly system. Assembly systems can be classified as (1) driven assembly systems and (2) self assembly systems or a (3) combination of both.

In driven assembly systems, the device is guided from their source and placed at their final destination while the entire process is monitored. The process is non-statistical and is mostly independent of the device and substrate properties. The surface mount technology (SMT) pioneered and developed by IBM since 1960's is an example of such a driven system. SMT is used extensively in the electronics industry for soldering components onto circuit boards [1]. Recent developments in SMT include the heterogeneous device assembly technique called "pick and place" method developed utilizing a robotic arm [2]. Components are manufactured separately and assembled in a serial manner. Due to the continuing miniaturization of the devices to be assembled, this method may become a bottleneck in terms of the precision.

An alternative approach to the driven assembly method is the self assembly technique. Nature leads the way for technology to follow self assembly techniques. From atomic level in crystal growth to tissue level in biological systems, self assembly is directed by energy and entropy considerations. Following Whitesides and Boncheva [3], "Self-assembly can be defined as a process by which preexisting components (separate or distinct parts of a disordered structure) autonomously organize into patterns or structures without human intervention [3]". Self Assembly system components can be self propelled or externally propelled. External propulsion systems may use mechanical agitation, gravity or electromagnetic fields to increase the chances of component interactions. Once the components are within interaction distance, they may

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assemble to form the desired final product. Self propelled systems on a macroscopic scale are mostly robotic systems in which components are powered to change position orientation or conformation and combine to form larger modules. A comprehensive list of natural and man-made self assembly systems are listed in [4]. Self assembly techniques developed since the 1980's include Fluidic Self Assembly (FSA) developed by Yeh and Smith [5] using the concept of assembly that uses gravity, fluid flow and shape recognition to drive the assembly of micro devices. A trapezoidal shape of devices is chosen to constrain the orientation of the devices moving under gravity. This process was used to successfully integrate GaAs light emitting diodes (LEDs) and other similar components onto silicon substrates. FSA is reported to have a yield ranging from 30% - 70% for a device size of around 20 urn [6] and a near 100% yield for devices of sizes in the range of 100 um to 1.5 mm [7]. Another approach is by the capillary force driven self assembly method pioneered by Whitesides' group at Harvard University. This method utilizes the lateral capillary forces to create adhesion between hydrophobic regions of the components to form two dimensional self assembled systems [8]. Many variations of FSA utilizing capillary forces have been successfully developed and incorporated into the microelectronics industry [9 - 12]. Other self assembly techniques involving the use of electric and magnetic fields have also been developed. Some of the notable assembly techniques are Magnetically Assisted Statistical Assembly by Fonstad at MIT [13] using micro magnets for um scale devices. Devices and surfaces (with recesses) were fabricated with shapes and sizes complementary to each other. A high coercivity permanent magnetic layer such as Co-Pt alloy was deposited and patterned onto the recess bed. The parts containing a layer of Ni fell into the recesses due to gravity and remained attached to the recesses when the "snow globe" apparatus was overturned.

Deterministic Approaches: Many of the self assembly processes depend, to a great extent, on the statistics of device position and movement under the influence of various forces. This is due to chaotic behavior of the interacting forces. Even in a controlled environment, the number of forces that arise due to system components' interaction among themselves and the environment is large enough to create a chaotic scenario. In many of the above mentioned self assembly processes, the number of devices far exceed the number of recesses available so that the integration statistics become favorable and a high yield rate is achieved. Ramadan's group at the University of Singapore has demonstrated a high yielding technique using nearly three times the number of devices to be placed for a 97 % yield in 5 minutes [14]. In order to make the small device assembly deterministic, one needs to eliminate a few forces and introduce constraints that would make the motion of the devices predictable to a very high degree of accuracy. At New Jersey Institute of Technology, Prof. Ravindra's group studies various deterministic approaches to the magnetic field assisted assembly [15], The substrate has recesses with a magnetic layer and shape of the recesses match the device shape and size. The devices also have a magnetic layer that responds to the field created by a moving magnet located below the substrate. This magnet guides the devices into the recesses where they get attached to the substrate. A detailed device-substrate force calculation has been presented by Rivero et. al. [16]. Miura's group has demonstrated controllable motion of the devices [17].

This work is related to the recent progress made by Prof. Ravindra's group in the field of magnetic field assisted assembly [18], Here the success of device assembly is always high provided there is no significant influence from the environment. Further, statistical analysis in such deterministic methods is employed to understand the failure rate of the system components rather than a success rate or yield of the device assembly itself. In this study, we use device and

652

substrate integration using magnetic fields generated by an array of solenoids. The solenoids guide the devices into their final position. The path taken by the devices on the substrate is predetermined. The device and solenoid geometry is such that the devices do not stray away from their predetermined path. Simple sensing techniques are incorporated to detect misfits and misalignments. Any error is detected right away and corrective measures are implemented without resorting to iterative techniques. The magnetic fields generated by various solenoid geometries is analyzed and the shape that provides a particular radial field strength for a given current is large enough to overcome the frictional forces and the transverse component helps in controlling frictional forces. Device geometry is designed with a minor modification to existing shapes that are in extant and the modification can be incorporated into a packaging process with little or no extra cost.

Designing the Magnetic Field Assisted Assembly system

Physics of Magnetic Field Assisted Assembly

Field Distribution: The magnetic field generated by a current carrying solenoid can be calculated using the Biot-Savart law given byB= L 0 M x r / r ' . The field due to a cylindrical coil of length 2b and n turns has a field distribution given by:

ß = 2B0 \dz[\— 2 aZ

2COS<* —d*) ( ')

. = 2*o f {]— ? -^cos* jj^dAdz (2)

and

where, Bp and Bz are the radial and transverse magnetic field components respectively [19]. B0 is the field along the axis of the solenoid given by B0 = ß0nl I An, / being the current passing through the circuit. When a magnetic dipole is placed in the magnetic field created by the coil, the interaction between the dipole and the field is given by z = ß x B, x being the torque acting on the magnet due to the transverse component of the field (eq. 2) and m, the magnet's magnetic dipole moment. Due to the finite size of the magnet, the solenoid field strength at the ends of the magnet will not be identical creating a net force on the magnet. The larger the dimensions of the magnet, the greater will be the difference between the fields at its ends. The force on the magnetic dipole due to a field gradient is given by F = V(// • B), which moves the magnet across the surface. From the force and torque equations, it is obvious that the components of the fields must be manipulated in order to optimize the motion of the magnets and decrease the deviation of the magnet's movement from a predetermined path. The geometry of the solenoids determines the field distribution. Since the objective is to make the magnets move on a horizontal surface in a controlled manner, the normal component of the magnetic field must be made to vary in a fairly linear manner. For a cylindrical solenoid, the transverse field varies as 1/r3, which implies that, even for a small r, the gradients are very steep. Such steep gradients accelerate the magnets making them deviate from a desired path. To prevent large accelerations, one needs to create a field that falls off slower than 1/r3. In order to obtain a linearly varying field, one needs to create

653

an appropriate geometry. The geometries of the cores used in the coils are modified to redirect the magnetic flux in such a way that the transverse components have a gentler gradient.

Forces and Torques: A magnet placed in a non uniform magnetic field experiences both a torque and a force. Further frictional forces exist between the magnet and the surface on which it rests. The net force acting on a magnet of mass M that is in contact with a horizontal surface of a coefficient of friction ur can be modeled as:

F = V(M-B)-F/ricllo„ (3)

F,=0,F.=t'.^,Fp = fi.^-ti,(Mg + F,) (4)

The transverse force also accounts for the increase in the friction between the two surfaces. Coating each magnet with a thin layer of poly tetrafluoroethane PTFE (Teflon®) or similar material will help to minimize the friction. The radial acceleration of a device is given by:

M M

The time dependence has to be explicitly included to account for the transient nature of the fields. The change in magnetic field intensity from zero to a desired value and back to zero depends on the time constant (L/R) of the solenoid circuit and the hysteresis of the core used which leads one to the time dependent acceleration aft) and radial path p(t) given by:

< < / ) = ^ ( / i . ^ ^ - / i / f ^ + ̂ ) > | ] ( l - « p ( - * / i ) ) - « ( 0 ) ( l - e x p ( - / a / i ) ) (6)

and

^ - M M g + * & il (5)

p(')-

(Mr 3Ä2(0) ( dB ,(0)

dz

f « ' A

(o) Z 2 ( l - e L) l2

— - ; - + — R2 2

(7)

R and L are the resistance and the inductances of die circuit. The force acting on the magnet does not increase monotonically with time as suggested by eq. (7) since the gradient of the field is position dependent. A similar approach can be used to calculate the net torque acting on a magnet of dipole moment u of length /, width w and mass M. It will undergo an angular acceleration aft) given by:

B W = I = ̂ m (8) a(t) helps to determine the angle turned in time t. Equations (1) through (8) describe the motion of the magnet that is located on top of a substrate. The substrate resides above an array of solenoids. It is to be noted that we neglect the velocity dependent terms of resistance due to the dry nature of friction throughout the setup.

Force calculations

Let c = (z2+a2+p2)/2ap and m = 2!(\ + c) = 4ap/z2 + (a + p)2. Theneq. (1), (2) can be written as

_ _ az f — c o s <pd<p m

654

■2Bn (lap)

, "r d^ "r — cos (t>d<t> "2 C a | (C - cos«*)"2 + P\ ic - cos«*)3'2 >

(10)

The fields can be expressed in terms of the Elliptic Integrals of the 1st (E) and 2nd (K) kind as follows [21]:

B„ = - 2 Ä , 1 — m p Aap 2 (1 — m) -) E(m) - Kim)

Bz = 2B0i-^y> [ pKim) + am-\2 ~ 7>^ Eim)

Aap \ 2(1 — m)

(11)

(12)

The Force acting on a bar magnet whose magnetic moment is aligned along z is given by eq. (4). Using the above form of the fields, we can calculate the force on a magnetic dipole. The radial fields of a solenoid and the forces on a test dipole due to a solenoid are shown in Figures lc and lb respectively. Figure la is the normalized radial component of the force acting on the device (magnets) on the substrate. As expected, the radial force goes to zero on the axis of the solenoid and the transverse component Fz becomes large. This is the component that holds the magnet at its desired position while further steps are taken to fix it during the assembly process.

Figure 1. (a) The normalized force F = Fr?2=a/Fr,z=o experienced by a small magnetic dipole when placed at a height z = a (the radius of the coil), (b) Shows the complete force map at z = 2a above the coil that produces a field (c).

The radial component of the force is given by:

F = uja-P 0

litp\ ((>-/>* ^ . 2 (1 + p)

/+(, + z 2 ) 2+ / ( . 1 + 2 l2) + ^ | . 1 . , 2 r 2 + : 4 j £ *P

(-((-1 + P ) : *+S})\I + P*+,2+P)[-I+S)K 4p

(-7 + 6P2+2(-3 + P1)Z2+Z*)E 4p

(! + />)

(1 + P) +■

(-i + PKS)K

2*((l-p) :

Ap

ill

Ul + />)2+z2 )

Ul+P)2+--2 / (14)

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Eq. (14) represents the force on a dipole due to a single coil of radius 1mm. An array of coils when turned on in a sequential manner can be used to drive a magnetic dipole from one point to another. The timing of the "switch on" of the coil is predetermined by solving for the time it takes to move across a given distance (eq.7) under the net horizontal force Fp (eq. 14). Using the above equations, the displacements of magnetic dipoles of different dimensions have been calculated. The results are summarized in Table 1.

Tablel.Simulation of force on devices of various sizes and die corresponding displacements using a current of 250 mA in a single layer coil of diameter 1 mm and 250 turns of 0.05 mm ( AWG 38) Cu wire. The coefficient of friction is taken to be 0.1 PTFE on polished silicon.

1

2

3

4

5

Device dimensions

(1 x w x h)

(mm)

1x0.6x0.3

2x1.25x0.5

3x1.875x0.5

4x2x0.5

M

(mg)

0.99

4.35

165

220

(Am2)

4.5» 10'5

3.2'W4

9»10""

1.5*10'3

Field Details at (p/a = 2) at t » t

Current

(A)

0.25

0.25

0.25

0.25

aB/Sp

(mT/mm)

1

1

1

'

dBJSp

(mT/mm)

1.5

1.5

1.5

1.5

(uk=0.1*)

uN

8.2

105

200

276

Net Force

uN

16

25

40

42

Choice of Core and coil shape. In order to create a force that varies in a way that the motion can be controlled efficiently and without jerks and jitters, one should develop a field gradient that is more or less constant since the field gradient determines the force. A few simple geometries of the coils and cores were studied and their fields are plotted in Figure 3.The fields and forces due to these coils are determined using finite element methods. This was accomplished using COMSOL 3.5 [19].The fields for various geometries of coils/cores are plotted in Figure 3. The least gradient is that of the cylindrical coil with a spherical core. This shape is found to produce the smallest variation of force with radial separation between the dipole and the solenoid.

Fig . 2. (a) Cylindrical core for a straight solenoid, (b) Conical coil with a conical core, (c) Spherical top core for a straight coil. The spherical part projects out. In all three cases, core radii are nearly equal to the inner radii of the coils to contain as much field as possible and prevent large leakage.

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Figure 3. Variation of the transverse (a) and the radial (b) components of B-field for various coil geometries show that the field gradient of a spherical top core has the least value(triangles) The top and bottom set of plots are transverse fields at z = 1 mm and 2 mm respectively above the top of the coil.

Designing the device carriage: The devices ride on a rectangular frame with magnets at the vertices. Their magnetization is perpendicular to the plane of the substrate. The dimensions of the device frames influence the turning moments for which they should be optimized. The magnets used in the device carriage should be such that the magnetization should not interfere with the device functions. To conform to this demand, a low coercivity (soft) material is chosen. The magnetization can be removed by flipping the polarity of the current in the solenoids to produce a coercive field for a sufficiently long time after the device is fixed in its final position. The shape of the magnets is chosen to be cylindrical to minimize regions of high local magnetizations that are associated with the sharp edges of cuboidal magnets [19]. Cylindrical magnets possess a uniform magnetization, which helps in eliminating any shaking while moving under the solenoid fields. Figure 4a shows a schematic of the device carriage with magnets attached. The diameter of the magnets is chosen to be at most a quarter of the shorter side of the device carriage and the height of the magnets is determined by the depth of the recesses on the substrate. The magnets are coated with a very thin layer of polymer (Teflon, for example) to decrease the resistive forces. Based on the theory discussed earlier, it is possible to assemble devices on a substrate with a high precision. A substrate with recesses will accept devices which

657

move into their respective positions, following a predetermined path; the motion and orientation is driven by the solenoid array. A sequential turning on and off of the solenoids will generate the field gradients which will guide the devices along the gradient. Thus the linear displacement can be achieved easily. Since the poles of the magnets used in the device are vertical while the device movement is horizontal, torque cannot be used to incorporate any angular displacement in the horizontal plane. The only way to achieve this is by using multiple magnets on the devices. Figure 4 shows a choice of arrangements of magnets to facilitate controlled angular displacement of a device. The shape of the device itself has lesser influence on the assembly process compared to the position of the magnets on the device. Multiple magnets on the device provide enough asymmetry to produce torque in the device. In Fig. 4b, a 90° turn is displayed. The number of steps can be reduced greatly by turning the fields simultaneously thereby reducing the operation time. In the following sections, each component of the assembly system is discussed in detail.

Figure 4 (a) Device housing with magnets attached, (b) Turning a device by 90° using solenoids in the vicinity of the magnets. The solenoids form a part of a larger array. The field direction and intensities are the following: Vertical (Horizontal) lines indicate intensity of the field directed out of (into) the paper. (1) A large current is passed through Bl to fix the device's leg on Bl. (2) Smaller currents are passed through Al , A2 and C2 to begin the turning. (3) Next ,B2,and C2 are also turned on . C2 and B2 are turned off. (4) Currents in Bl and Cl are increased. (5) All currents are turned off when the device reaches its final position.

Designing the device carriage: The devices ride on a rectangular frame with magnets at the vertices. Their magnetization is perpendicular to the plane of the substrate. The dimensions of the device frames influence the turning moments for which they should be optimized. The magnets used in the device carriage should be such that the magnetization should not interfere with the device functions. To conform to this demand, a low coercivity (soft) material is chosen.

658

The magnetization can be removed by flipping the polarity of the current in the solenoids to produce a coercive field for a sufficiently long time after the device is fixed in its final position. The shape of the magnets is chosen to be cylindrical to minimize regions of high local magnetizations that are associated with the sharp edges of cuboidal magnets [19]. Cylindrical magnets possess a uniform magnetization, which helps in eliminating any shaking while moving under the solenoid fields. Fig. 4.a shows a schematic of the device carriage with magnets attached. The diameter of the magnets is chosen to be at most a quarter of the shorter side of the device carriage and the height of the magnets is determined by the depth of the recesses on the substrate. The magnets are coated with a very thin layer of polymer (Teflon, for example) to decrease the resistive forces.

Scaling issues: One important aspect of this work is to clearly define a limit on the size of the devices that can be assembled. While it is fairly easy to build millimeter sized coils and devices, it is a big technological challenge to build a system that can assemble micrometer size devices. For a given size of the coil, the size of the device has a lower limit: The separation between the magnets attached to the device must be larger than the diameter of the coil. Only under this condition, the device can be made to turn (as shown in Fig. 4b) with minimal uncertainty. Further, scaling offerees with size has to be quantified. The field produced by tiny coils (10-50 urn) can be controlled with almost arbitrary precision due to the availability of modern deposition techniques and power supplies [19], but creating a micro magnet of desired magnetization orientation is a challenge.

Sophisticated techniques have been employed to create NiFeCo micro magnets of patterned magnetization for the data storage industry [20]. Such techniques may be utilized to produce micro magnets with uniform magnetization. Currently, various techniques of fabricating magnets and coils are being explored.

If the requirement is for a large device assembly, then instead of one solenoid per leg, a bunch of solenoids can be used to move the device. Therefore the same system can be used to assemble devices that scale from the solenoid diameter to any size larger than the solenoid diameter.

Stability of the device is also affected due to the small size. In order for the device to slide over the substrate, the distribution of mass and the orientation of the dipole moment become critical, since both are parallel to each other. When the device is moving under the influence of fields, the dipole moment causes the device to lean forward due to the radial force. As it is seen from the force and the field plots, the gradients are large in the vicinity of the coils, which will give rise to large radial accelerations. If the switching times of the coil is not matched with the position of the devices, large accelerations will cause the device to topple or flip. To counter this effect, the devices are required to be housed in a carriage whose in plane dimensions are designed to minimize such toppling due to the torques acting on the devices.

Conclusions

In this paper, the concept of magnetic field assisted assembly is discussed. The physics of device movement is presented. Forces acting on the devices of submillimeter dimensions are calculated. The motion of the device in a single array of solenoids is simulated using finite element methods and the results are presented. The aim of this study is to present a viable solution for heterogeneous device assembly.

659

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[I] Leonard J. Poch "High speed packaging of miniaturized circuit elements", United States Patent Number 3312878 4/4/1967

[2] K. F. Böhringer, R. S. Fearing, and K. Y. Goldberg, "Microassembly," in Handbook of

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[3] G.M.Whitesides and M. Boncheva, B "Beyond molecules: Self-assembly of mesoscopic and

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[4] Roderich Grob, Marco Dorigo, "Self-Assembly at the Macroscopic Scale" Proceedings of the

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[5] J. S. Smith, "Fluidic Self-Assembly of Microstructures and its Application to the Integration

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[6] H. J. Yeh and J. S. Smith, "Fluidic self-assembly for the integration of GaAs light emitting

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[8] N. Bowden, A. Terfort, J. Carbeck, and G. M. Whitesides, "Self-Assembly of Mesoscale

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[10] J. Tien, T. L. Breen, and G. M. Whitesides, "Crystallization of Millimeter-Scale Objects

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12671, 1998.

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[13] C. G. Fonstad, "Magnetically-Assisted Statistical Assembly - a new heterogeneous

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assembly using an external magnetic array "Applied Physics Letters Vol. 90 Issue 17,2007.

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[15] S. Shet, V.R. Mehta, A.T. Fiory, M.P. Lepselter, and N.M. Ravindra, "The Magnetic Field-

Assisted Assembly of Nanoscale Semiconductor Devices: A New Technique" JOM 56, 32

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[16] Rene Rivero, Sudhakar Shet, Michael R. Booty, Anthony T. Fiory, and Nuggehalli

Ravindra, "Modeling of Magnetic-Field-Assisted Assembly of Semiconductor Devices" Journal

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"Micromanipulation Using Magnetic Field" IEEE international Conference on Robotics and

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7,737,515 B2.

[19] COMSOL Multiphysics Version 3.5, COMSOL Inc.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ORGANIC THERMAL MODE PHOTORESISTS FOR APPLICATIONS IN NANO-LITHOGRAPHY

Hsiu-Wen Wu1, Ming Chia Li2 ,Chin-Tien Yang3, Chung-Ta Cheng2, Shuen-Chen

Chen2 and Der-Ray Huang*1

1 Dept. of Opto-Electronics, National Dong Hwa University, Hualien 970, Taiwan 2Electronics and Optoelectronics Research Laboratories, ITRI, Hsinchu 310, Taiwan

3Nanotechnology Research center, ITRI, Hsinchu 310, Taiwan

Keywords: Thermal photoresists, Submicron pattern, nano-lithography

Abstract

A new technology called thermal mode lithography for developing new thermal mode

photoresists with temperature sensitive characteristics has been investigated. In this

study, we synthesized four types of polymethine compounds that can be applied to

thermal mode photoresists. The photo UV absorption and thermal properties are the

important characteristics of thermal mode photoresists. The decomposition and

gasification mechanisms of thermal photoresists can be analyzed by thermal

properties. The surface morphology of nano-patterns can be controlled by using

different laser power exposure on thermal photoresists. A cost effective way to

achieve submicron or nano-structure patterns is possible.

Introduction

For applications in nano-technology, nano-lithography is the most important process

to achieve it. Currently, some nano-lithographic technologies such as using

electron-beam (e-beam), atomic force microscopy (AFM), X-rays etc., are possible to

write nano-scale pits. However, these processes for mass production with high

efficiency are still impossible. Right now, the commercial laser beam lithographic

system is popular and is cost effective for submicron-scale industry applications.

But the diffraction limit of laser beams will limit it to achieve nano-scale pits even by

using deep ultraviolet (DUV) laser with short wavelength 248nm/257nm/266 nm and

high numerical aperture (NA) of 0.9. In order to solve the diffraction limit issue, we

have investigated to develop thermal mode photo-resistors to control pit size. A new

technology called thermal mode lithography aims at developing a new photoresist

with temperature sensitive characteristics that can breakthrough optical diffraction

limit [1]. The photo UV absorption and thermal properties are important

characteristics of thermal mode photoresists [2]. It can control the submicron pattern

size and resolution. Most thermal mode photoresists are manufactured by using

663

inorganic metal oxides that have pollution problems in semiconductor processes. In

this study, we have synthesized organic compounds that have applications in thermal

mode photoresists and use thermal lithography technology to prepare submicron

structures.

Experimental Details

In this study, we synthesized four types of polymethine compounds (Al, A2, A3,

A4) that can be applied to thermal mode photoresists. The polymethine powder is

dissolved in a solvent to prepare the solutions at 2wt% concentration. Then the

solutions are spin-coated on sapphire and silicon substrates to form photoresist layer

at a rotation speed of 2500 rpm. Following this step, the thin film samples were dried

to remove the solvent at 90°C for lmin and 55°C for 40min, respectively. The

thickness of the thin film layer are controlled between 80~200nm. A new laser

opto-mechanical system with high precision R-0 rotation has been developed by ITRI

EOL. A single wavelength laser diode was used as the exposure source that can

directly focus on a rotation stage [3] (Fig.l). The operating condition of equipment

was monitored by using blue laser wavelength of 405nm, NA 0.85 and stage rotation

speed of 2.4 m/s. The characteristics of UV absorption and thermal properties were

monitored by using of UV-Vis, TGA and DSC. The morphology of the submicron

patterns of thermal photoresists on substrates was measured by using AFM.

Fig. 1 ITRI EOL opto-mechanical system for direct laser writing

664

Results and Discussion

UV absorption and thermal properties are the important parameters to analyze

material degradation mechanisms of thermal mode photoresists. Thermal mode

photoresists can be excited by blue laser 405nm, and then, by gasification,

submicron patterns can be formed. The degradation curves by TGA analysis are

shown in Fig.2. The Al compound degradation curve is observed by the first-order

decay and the second-order decay, while other compounds are just observed by the

first-order decay. Fig. 3 shows the UV-VIS absorption curves of four type A series

polymethine compounds. The absorption peaks of four samples Al, A2, A3, A4 are

450nm, 450nm, 464nm, 489 nm respectively. From the results of UV absorption

curves, it can be noticed that A4 compound needs bigger laser power exposure. All

thermal analysis data are shown in Table 1. The decomposition and gasification

mechanism of thermal photoresists are analyzed by thermal properties such as

melting temperature(Tm) and cracking temperature (Td). Results of the reaction of

exposure of the four types of polymethine compounds (A series) are exothermic and

have high enthalpy. The A4 compound has small pit size because of its high

enthalpy and decomposition area. The morphology of thermal mode photoresists on

sapphire substrates can be controlled by the power of blue laser.

Fig.2. Degradation curves of A series - DTG curves

665

Fig.3. Absorption curves of A series compounds (solvent MeOH)

Table 1. Thermal properties of A series photoresists

Compound

Tm(°C)

AHm(KJ/mole)

Td(TC)

DTG(°C)

Al

183.68

-63.145

199.53

220.53

295.59

A2

173

-91.129

236.53

265.5

A3

265.52

-121.4

270.23

336.87

A4

252.94

-61.51

298.27

342.39

Fig.4 shows the submicron patterns of thermal mode photo-resists on sapphire

substrates that were observed by AFM. The laser power used for samples Al, A2, A3

and A4 are 3mW, 4mW, 6mW and 7mW respectively. Comparing the morphology of

different thermal mode photoresists, Al has good surface morphology but pit size is

too big, A2 and A3 show high thermal conductivity with low enthalpy, and A4 has

good performance of pit size and depth. The pit size of sample A4 is close to optical

diffraction limit. With some modification of process, it is possible to breakthrough

optical diffraction limit and good surface morphology.

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Fig.4. The submicron patterns of photoresists on sapphire substrates (a) Al : pit 449nm

(b) A2: pit 722nm (c) A3: pit 624nm (d) A4: pit 321nm

Conclusions

An innovative technology to produce large scale and well ordered submicron

patterns on sapphire substrates has been developed by using low-cost direct laser

writing process. The patterns of organic photoresists are obtained by laser diode and

gasification. This mechanism with nonpolluting semiconductor processes can lead to

submicron or nano-structure patterns with fewer steps. The morphology of thermal

mode photoresists on sapphire substrates can be controlled by the power of blue laser.

It is possible to breakthrough optical diffraction limit and obtain good surface

morphology. This will be very useful for applications in some new nano-lithographic

processes.

Acknowledgements

The work was supported by the National Science Council, Taiwan, under Contract No.

NSC 100-222l-E-259-016. The authors would like to thank ITRl for providing

equipment and supporting the research.

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References

1. M. Kuwahara,*, J.H. Kim, J. Tominaga, "Dot formation with 170-nm dimensions

using a thermal lithography technique," Microelectronic Engineering, 67(2003),

651-656.

2. Y Usami et al, "405nm Laser Thermal Lithography of 40 nm Pattern Using Super

Resolution Organic Resist Material," Applied Physics Express, 12(2009) 126502.

3. C T Yang, "Single Wavelength Blue-Laser Optical Head-Like Opto-Mechanical

System for Turntable Thermal Mode Lithography and Stamper Fabrication", IEEE

Transactions on Magnetics, 47(2011), 3.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

NITROGEN DOPED ZnO (ZnO:N) THIN FILMS DEPOSITED BY REACTIVE RF MAGNETRON SPUTTERING FOR PEC

APPLICATIONS

Sudhakar Shet,1,2' Kwang-Soon Ahn,3 Nuggehalli Ravindra,2 Yanfa Yan,1 and Mowafak Al-Jassim1

1 National Renewable Energy Laboratory, Golden, CO 80401 USA 2 New Jersey Institute of Technology, Newark, NJ 07102 USA

3School of Display and Chemical Engineering, Yeungnam University, Dae-dong, Kyungsan 712-749, South Korea

Keywords: N-doping, photoelectrochemical, thin film

Abstract

ZnO:N films were deposited by reactive RF magnetron sputtering on F-doped tin oxide coated glass substrates in mixed N2 and O2 gas ambient. Their PEC properties were measured and compared with those of as-deposited and annealed ZnO films. The ZnO:N films exhibit photoresponse in the visible-light region, yielding higher total photocurrents than ZnO thin films. ZnO:N thin films with reduced bandgaps were synthesized by reactive RF magnetron sputtering using ZnO target at 100°C followed by post deposition annealing at 500°C in air for 2 h. ZnO:N thin films showed enhanced N incorporation and shift of the optical absorption into the visible light regions. As a result, ZnO:N films showed improved PEC response, compared to ZnO thin films

Introduction

Transition metal oxide-based photoelectrochemical (PEC) splitting of water has attracted wide interest since photo-induced decomposition of water on Ti02 electrodes was discovered [1]. To date, most investigations have focused on Ti02 [2-7]. The drawback of a PEC system using TiÛ2 is that it can only absorb ultraviolet (UV) light due to its large bandgap of 3.0-3.2 eV. Therefore, it is necessary to search for new metal oxides that can potentially absorb visible light. ZnO has similar bandgap and band-edge positions compared to TiÛ2 [2]. Furthermore, ZnO has a direct bandgap and higher electron mobility than Ti02 [8]. Thus, ZnO could also be a potential candidate for photoelectrochemical applications [9]. However, like Ti02, the bandgap of ZnO (-3.3 eV) is too large to effectively use visible light. Hence, it is critical to reduce its bandgap to achieve a higher absorption coefficient.

Impurity doping in the photoactive metal oxides has been known to shift light absorption to longer wavelengths. Asahi et al. [4] recently reported that nitrogen doping was one of the most effective methods for bandgap narrowing, which led to improved photoresponse in the long-wavelength region. Similarly, N incorporation in ZnO is demonstrated to reduce the optical bandgap, leading to absorption in the long-wavelength regions [10-

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11]. However, to date, the PEC properties of N-incorporated ZnO (ZnO:N) have scarcely been studied.

In this paper, we measured the photoelectrochemical properties of ZnO:N thin films prepared by reactive RF magnetron sputtering in mixed N2 and 0 2 gas ambient and compared them with those of as-deposited and annealed pure ZnO films. Also effect of RF power on the ZnO:N thin films and their PEC properties were studied. We find that N incorporation can red-shift the onset of photoresponse, and thus yield higher total photocurrents than pure ZnO films. Our results provide insights on the use of ZnO:N-based PEC systems.

Experimental Details

Two sets of samples were prepared. In the first set, Zn target was used and in the second set, ZnO target was used. For the first set, the ZnO:N films were deposited by reactive RF sputtering of a zinc metal target using an oxygen/nitrogen gas mixture. F-doped Sn02-coated transparent glasses (20-23 fi/D) were used as substrates. The distance between the target and substrate was 8 cm. The base pressure was below 5><10"6 torr and the working pressure was 2x 10"2 torr. The chamber ambient was mixed N2 and 0 2 with an oxygen gas ratio 02/(N2+02) = 5%. This low 0 2 gas ratio was found to be necessary because the chemical activity of 0 2 is much higher than that of N2. Prior to sputtering, a pre-sputtering process was performed for 30 min to eliminate any contaminants from the target. Sputtering was then conducted at a RF power of 80 W at room temperature. For comparison, ZnO films were deposited at a RF power of 80 W in pure 0 2 gas ambient. All the sputtered samples were controlled to have a similar film thickness of 500 nm as measured by stylus profilometry. For the second set, the ZnO:N thin films were deposited by reactive rf magnetron sputtering of a zinc oxide target using an oxygen/nitrogen gas mixture. F-doped Sn02 (FTO, 20-23 fi/Q)-coated transparent glasses were used as substrates. Sputtering was then conducted with various RF powers in the range of 100 to 300 W at 100°C substrate temperature. For comparison, ZnO films were deposited at an RF power of 200 W in pure Ar ambient. All samples were controlled to have similar film thickness of about 1000 nm as measured by stylus profilometry.

For structural characterization, X-ray measurements were performed using an X-ray diffractometer (XGEN-4000, SCINTAG Inc.), operated with a Cu Ka radiation source at 45 kV and 37 mA. The ultraviolet (UV)-visible (Vis) absorption spectra of the samples were measured by an n&k analyzer 1280 (n&k Technology, Inc.). To evaluate the N chemical states and their concentrations in the ZnO:N films, X-ray photoelectron spectroscopy (XPS) was performed. Monochromatic Al Ka radiation was used for all data sets, and the analyzer was set to 59-eV pass energy. Argon ion sputtering (3 keV, 0.8 nAmm"2, 120 s) was used to clean samples prior to analysis. Photoelectrochemical measurements were performed in a three-electrode cell with a flat quartz-glass window to facilitate illumination to the photoelectrode surface [12-17]. The sputter-deposited films were used as the working electrodes with an active surface area of-1.2 cm2. Pt mesh and a Ag/AgCl electrode were used as counter and reference electrodes, respectively. A 0.5-M Na2SÛ4 aqueous solution with pH of 6.8 was used as the electrolyte for PEC measurements [18-21]. The photoelectrochemical properties of the samples were measured using a fiber-optic illuminator (150-W tungsten-halogen lamp) with a

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UV/infrared (IR) filter. AC impedance measurements were carried out with a Solartron 1255 frequency response analyzer using the above three-electrode cells. The series capacitor-resistor circuit model was used for Mott-Schottky plots [22].

Results and Discussion

X-ray diffraction curves of the as-grown and annealed ZnO and ZnO:N films are shown in Fig. 1. The crystallite sizes estimated according to the Scherrer equation for the as-deposited ZnO film exhibits nanocrystalline features with crystallite sizes below 10 nm. However, the as- deposited ZnO:N films showed a sharper peak, indicating larger crystallite sizes of about 24 nm. Deposition rate analyses show that the growth rate was 3.7 nm/min for pure ZnO, but 9 nm/min for ZnO:N films. Because the crystallinity of the film is closely related to its PEC property, the pure ZnO film was annealed for 1 h at 500°C in air to achieve similar crystallinity to that of ZnO:N. After annealing, the crystallite sizes were shown to have increased to about 22 nm, close to that of the ZnO:N films, as shown in Fig. 1.

Figure 1. X-ray diffraction curves of the as- deposited ZnO, ZnO:N, and annealed ZnO films.

Optical absorption spectra of the as- deposited and annealed ZnO films and the as-deposited ZnO:N films is shown in Fig. 2.

Figure 2. UV-Vis optical absorption spectra of the as- deposited and annealed ZnO and ZnO:N films.

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The as- deposited and annealed ZnO films, due to their wide bandgap, showed similar optical absorption spectra and could absorb only light with wavelengths below 450 nm. However, the ZnO:N films could absorb lower-energy photons, up to 700 nm, indicating that the bandgap was narrowed by N incorporation in ZnO [10-11].

Figure 3 shows Mott-Schottky plots of the as- deposited and annealed ZnO films and the ZnO:N films. The positive slopes indicate that all of the samples are n-type semiconductors. The origins of the «-type conductivity in the as- deposited and annealed ZnO films are likely due to native defects such as oxygen vacancies (V0) or/and Zn interstitials (Zni). It is known that a N occupying an O site (Nc) is an acceptor, which should result in p-type ZnO:N. However, Fig. 3 shows that the ZnO:N film was actually «-type. It was reported that the ZnO films deposited in N2 plasma are usually «-type, due to substitutional N2 molecules at an O site that act as shallow double-donors [23-24].

Figure 3. Mott-Schottky plots of the as- deposited and annealed ZnO and ZnO:N films.

The donor densities of the as- deposited and annealed ZnO films and the ZnO:N films are calculated to be 4.6xl016, 1.8*1016, and 3.8><1017/cm3, respectively. The annealed ZnO films exhibit lower donor density than the as- deposited ZnO films. This may be due to the annealing being conducted in air, which could reduce the amount of Zn interstitials and O vacancies. The ZnO:N films exhibit a higher donor density than the as- deposited ZnO films due to N2 incorporation and/or native defects such as Zn interstitials and 0 vacancies. The flat-band potentials (En,) of these films are obtained by the x-intercept of their Mott-Schottky plots. The as- deposited ZnO films and the ZnO:N films have similar Efb (~ -0.49 V vs. Ag/AgCl). However, the annealed ZnO films have an En, value of-0.38 V, which is about 0.1 V higher than the as- deposited ZnO films.

Photocurrent-voltage curves for the as- deposited and annealed ZnO films and the ZnO:N films under illumination with the UV/IR filter and dark currents are shown in Fig. 4. The anodic photocurrents of all the samples increase with increasing potentials, which is again indicative of «-type semiconductor (as corroborated in Fig. 3). The photocurrent onset potentials of the samples were about 0.2 to 0.4 V anodic from their Mott-Schottky-determined flat-band potentials. This overpotential is due to slow charge transfer and surface recombination at the semiconductor/solution interface.

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Figure 4. Photocurrent-voltage curves for the as- deposited and annealed ZnO and the ZnO:N films under light illumination) with the UV/IR filter and under dark condition,

indicated by the circle

Because the dark anodic currents begin to appear at -1.1 V, the maximum photocurrents of the three samples are compared at this potential. No limiting photocurrents are observed before the appearance of the dark anodic currents at 1.1 V. This phenomenon appears to be related to the small width of the space-charge layer of these samples. From the data of the Mott-Schottky plots in Fig. 5, the lva lues of the samples are calculated. At a potential of 1.1 V, the W values are about 181, 281, and 62 nm for the as- deposited and annealed ZnO films and the ZnO:N films, respectively. These W values at 1.1 V are much smaller than the film thickness of 500 nm, resulting in no limiting photocurrents. The photocurrent of the annealed ZnO films at 1.1 V is larger than that of the as-deposited ZnO films most likely due to both better crystallinity and larger value of W of the annealed ZnO films. It should be noted that the photocurrent of the ZnO:N films at 1.1 V is larger than that of the annealed ZnO films, despite its much smaller W value. Even with similar crystallinity (as shown in Fig. 1) and smaller W value, the larger photocurrent of the ZnO:N films indicates that the photocurrent is affected more significantly by the optical properties of the ZnO and ZnO:N films.

We will now discuss for the second set of samples. Figure 5 shows the X-ray diffraction curves of ZnO, and ZnO:N films deposited at various RF powers in mixed N2 and O2 ambient with 0 2 mass flow rate of 10%. It is seen that the ZnO films exhibit poor crystallinity, due to the low-temperature sputtering process. The ZnO:N films deposited at 100 W showed better crystallinity than the pure ZnO films, despite faster deposition rate. For ZnO deposition, the ambient is pure Ar gas. For ZnO:N deposition, the ambient is mainly N2 with only 10% 02 . When the RF power was increased to 200 W and 300W, the crystallinity again became poor. The ZnOfN films deposited at higher RF power grew faster compared to the deposition rate observed during 100 W deposition. The concentrations of N in 100W, 200W, and 300W samples were about 1.89 at.%, 1.98 at.%, and 2.45 at.%, respectively, as determined by X-ray photoelectron spectroscopy. It is known that a high concentration of dopant can deteriorate crystal structure. However, poor crystallinity of the ZnO:N film deposited at 200W and 300W may be caused mainly

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by the high deposition rate and to lesser extent by N incorporation in the films. Crystallite sizes were about 21, 42, 34, and 35 nm for the ZnO, ZnO:N(100W), ZnO:N(200W), and ZnO:N(300W) respectively, which were estimated by applying the Debye-Scherrer equation to XRD data.

Figure 5. X-ray diffraction curves of ZnO, and ZnO:N films deposited at different RF powers in mixed N2 and 0 2 ambient with 0 2 mass flow rate of 10%.

The ZnO films showed optical absorption spectra (not shown here) and could absorb only light with wavelengths below 450 nm, due to their wide bandgap. However, the ZnO:N film could absorb lower-energy photons, up to 1000 nm, indicating the N incorporation in ZnO. The absorption spectra show that the light absorption is successfully shifted into visible regions. The optical bandgaps of the films were determined by extrapolating the linear portion of each curve. The bandgap of the ZnO film is 3.26 eV, which is consistent with the results reported elsewhere. The direct optical bandgaps measured for ZnO:N films deposited at substrate temperature of 100°C, for 100 to 300W gradually decreased from 3.2 to 3.15 eV.

Mott-Schottky plots of ZnO and ZnO:N thin film samples exhibited positive slopes, indicating n-type semiconductors. Our previous studies reported that ZnO:N films deposited under a N2/02 plasma showed n-type behaviors, due to substitutional N2

molecules which act as shallow double-donors. Figure 6 shows the photocurrent-voltage curves of ZnO and ZnO:N thin films, under illumination with the UV/IR filter. It shows clearly that the ZnO:N films exhibited enhanced photocurrents, compared to the ZnO:N films. At a potential of 1.2 V, the photocurrents were 6.89, 21.14, 22.58, and 26.24 uAcm"2 for the ZnO, 100W, 200W, and 300W ZnO:N films, respectively. To investigate the photoresponses in the long-wavelength region, a green color filter (wavelength: 538.33 nm; FWHM: 77.478 nm) was used in combination with the UV/IR filter (not shown here).

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Figure 6. The photocurrent-voltage curves of ZnO, and ZnO:N films deposited at different RF powers under illumination with the UV/IR filter.

The ZnO films exhibited no clear photoresponse, due to its wide bandgap. The N doped ZnO:N films exhibited enhanced photocurrent than the ZnO films, despite much less light absorption. It indicates the high recombination rate of the photogenerated electrons and holes in the ZnO films, due to its inferior crystallinity. On the other hand, the N doped ZnO:N films exhibited increased crystallinity, reduced bandgap, which lead to enhanced photocurrent than the ZnO films. The results demonstrate clearly that the N incorporation in ZnO can shift the optical absorption into the visible light region, yet with an enhanced crystallinity and PEC response.

Conclusions

The photoelectrochemical properties of nitrogen-incorporated ZnO (ZnO:N) films prepared by reactive RF magnetron sputtering were measured and compared with those of pure ZnO films. We find that the nitrogen incorporation narrows the bandgap of ZnO and shifts the optical absorption into the visible regions. ZnO:N films exhibited greatly enhanced crystallimty compared to ZnO films. We further find that the ZnO:N films provide considerable photoresponse in the long-wavelength regions. As a result, the ZnO:N films exhibit higher photocurrents than pure ZnO films.

Acknowledgements

This work was supported by the U.S. Department of Energy.

References

[1] K. Honda and A. Fujishima, Nature (London) 238, (1972), p. 37. [2] M. Grätzel, Nature 414, (2001), p. 338. [3] T. Bak, J. Nowotny, M. Rekas, and C.C. Sorrell, InternationalJ. of Hydrogen Energy

27, (2002), p. 991.

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[4] R. Asahi, T. Morikawa, T. Ohwaki, K. Aoki, and Y. Taga, Science 293, (2001), p. 269.

[5] S.U.M. Khan, M. Al-Shahry, and W.B. Ingler Jr., Science 297, (2002), p.2243. [6] A. Ghicov, H. Tsuchiya, J.M. Macak, and P. Schmuki, Phys. Stat. Sol. (a) 203,

(2006), p. R28. [7] J. Yuan, M. Chen, J. Shi, and W. Shangguan, International J. of Hydrogen Energy

31, (2006), p. 1326. [8]. K. Kakiuchi, E. Hosono, and S. Fujihara, J. Photochem. & Photobiol. A: Chem. 179,

(2006), p. 81. [9] T.F. Jaramillo, S.H. Baeck, A. Kleiman-Shwarsctein, and E.W. McFarland,

Macromol. Rapid Comm. 25, (2004), p. 297. [10] M. Joseph, H. Tabata, and T. Kawai, Jpn. J. Appl. Phys. 38, (1999), p. L1205. [ 11 ] M. Futsuhara, K. Yoshioka, and O. Takai, Thin Solid Films 317, ( 1998), p.322. [12] K. -S . Ahn, S. Shet, T. Deutsch, C. S. Jiang, Y. Yan, M. Al-Jassim, and J. Turner, J.

Power Source, 176, (2008), p. 387. [13] S. Shet, K.-S. Ahn, T. Deutsch, H. Wang, N. Ravindra, Y. Yan, J. Turner, M. Al-

Jassim, J. Power Sources 195, (2010), p. 5801. [14] K.-S. Ahn, Y. Yan, S. Shet, K. Jones, T. Deutsch, J. Turner, M. Al-Jassim, Appl.

Phys. Lett. 93, (2008),. p. 163117. [15] S. Shet, K. -S . Ahn, N. Ravindra, Y. Yan, J. Turner, M. Al-Jassim, J. Materials 62,

(2010), p. 25, [16] S. Shet, K. -S . Ahn, T. Deutsch, H. Wang, N. Ravindra, Y. Yan, J. Turner, M. Al-

Jassim, J. Mater. Research 25, Doi: 10.1557/JMR.2010.0017, (2010), p. 69. [17] K.-S. Ahn, Y. Yan, S. Shet, T. Deutsch, J. Turner, and M. Al-Jassim, Appl.

Phys.Lett. 91, (2007), p. 231909,. [18] S. Shet, K. -S . Ahn, H. Wang, N. Ravindra, Y. Yan, J. Turner, M. Al-Jassim, J.

Mater. Science DOI 10.1007/sl0853-010-4561-x, (2010). [19] S. Shet, K. -S . Ahn, Y. Yan, T. Deutsch, K. M. Chrusrowski, J. Turner, M. Al-

Jassim, andN. Ravindra,/. Appl Phys. 103, (2008) p. 073504. [20] K.-S. Ahn, Y. Yan, M.-S. Kang, J.-Y. Kim, S. Shet, H. Wang, J. Turner, and M. Al-

Jassim, Appl. Phys. Lett. 95, (2009), p. 022116. [21] S. Shet, K. -S . Ahn, N. Ravindra, Y. Yan, J. Turner, M. Al-Jassim, J. Materials 62,

(2010), p. 25. [22] J. Akikusa and S.U.M. Khan, Inter. J. Hydrogen Energy 27, (2002), p. 863. [23] CL. Perkins, S.H. Lee, X. Li, S.E. Asher, and T.J. Coutts, J. Appl. Phys. 97,(2005),

p. 034907 [24]. Y. Yan, S.B. Zhang, and S.T. Pantelides, Phys. Rev. Lett. 86, (2001), p. 5723.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

SPIN-COATED ERBIUM-DOPED SILICA SOL-GEL FILMS ON SILICON

S. Abedrabbo123, B. Lahlouh3, S. Shet4, A. T. Fiory1, and N. M. Ravindra'

'Department of Physics, New Jersey Institute of Technology; Newark, NJ 07901, USA department of Physics and Engineering Physics, Stevens Institute of Technology;

Hoboken NJ 07030 USA 'Department of Physics, University of Jordan; Amman 11942, Jordan

"National Renewable Energy Laboratory; Golden CO 80401, USA

Keywords: Photoluminescence; Thin films; Silicon; Sol-gel processing

Abstract

This work reports optical functionality contained in, as well as and produced by, thin film coatings. A sol-gel process, formulated with precursor active ingredients of erbium oxide and tetraethylorthosilicate (TEOS), was used for spin-coating thin (-130 nm) erbium-doped (-6 at. %) silica films on single-crystal silicon. Annealed films produce infrared emission in the 1.5-um band from erbium ions in the film, as well as greatly enhancing (-100X) band-gap emission from the underlying silicon. The distinctly different mechanisms for the two modes of optical activities are interpreted in terms of optical emission theory and modeling; prospects for opto-electronic applications are discussed.

1. Introduction

Spin-coating of sol-gel films is widely applied as a cost-effective process for depositing nominally pure and impurity-doped thin silica films on substrates (see e.g. [1]). Of interest are optically active coatings suitable for silicon-integrated optoelectronics. The present method entails depositing erbium-doped silica films on silicon wafer material that are densified by thermal annealing. With suitable selection of thermal annealing temperature (Ta), photoluminescence measurements show strong emission in the 1.5-um wavelength band from Er+3 centers in silica (Stark-split intra-4f 4I13/2 - 4I15/2 transitions of Er+3 [2]) for T„ « 850 °C [3]; additionally, near band-gap emission at 1.16 (im from the silicon becomes greatly enhanced for Ta » 700 °C [4]. Section 2 describes experimental procedures and reports analysis of photoluminescence .from Er+3 in the sol-gel material. Section 3 presents analysis of photoluminescence from the silicon. General conclusions are presented in Section 4.

2. Optical Emission from Erbium

Optical telecommunication networks use rare-earth doped optical materials for a variety of applications, e.g. fiber and waveguide amplifiers, waveguides by index of refraction modification, infrared light sources, integrated optical devices, displays and lasers [2,5-8], and silica doped with erbium is a widely used optical material. Although Si02:Er can be prepared in film form by a variety of methods, spin-coating of sol-gels acquired special interest because of low deposition temperature and low overall cost [6,7,9-13]. For this work, a sol-gel technique was developed using En03 as the Er dopant source and introduced at high concentration (6 at.%) in the SiÛ2:Er film. The process yields optically active erbium-doped silica films at

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moderate annealing temperatures (T„ ~850 °C). An example of a suitable optoelectronic application is in fabrication of planar erbium-doped waveguide amplifiers (EDWAs).

Erbium Sol-Gel Processing

The erbium-doped sol was produced by hydrolysis of TEOS (tetraethylorthosilicate, Si(OC2Hs)4) in a solution containing erbium oxide and then annealing the gel. The relatively high Er concentration of 6 at.% is expected to form stable Er-0:Si-0 structures, owing to the strong Er-O bond (6.3 eV [14]) and the role of Er as a network element in glasses [15,16]. Special considerations of sol-gel processes concern quenching of the 1.5-um emission by resonant energy transfer to vibrations of residual OH [1,17] and dependence of external emission on Er concentration [10,11]. Much residual OH can be removed with high-temperature annealing (e.g. Ta ~ 900 - 1100 °C [8,18]); however, Ta < 1000 °C is preferable for integrated silicon-based optoelectronics and glass reflow processing [19]. In the process developed for the present work, T„ ~ 850 °C turns out to be sufficient for removing the effects of residual OH as well as reducing surface pores area to maintain low level of OH, thereby avoiding high-temperature processing.

Experimental Procedure

The starting solution for the sol-gel process contained 0.5 g Er2Û3 powder mixed into a solution of 4 ml ethanol, 4 ml acetic acid, and 1.6 ml deionized water mat was stirred at 45 °C for 3 hours. The sol was then prepared by the addition of 2 ml TEOS and stirring for 10 min at 80 CC. Hydrolysis with high water/TEOS molar ratio R (here, R = 10) with acid catalysis produces branched Si-O polymerization in the sol and leads to dense films [20]; the trade-off is an increased hydroxyl content to be removed by thermal annealing [13]. Following this step, the solution was passed through a syringe filter with 0.45-um pore size and spun coated on 2.5-cm pieces of cleaned Si (100) wafer substrates rotating at 1200 rpm for 30 seconds. The resulting gel films were then oven dried in air at 120 °C for 30 minutes. Post-deposition thermal treatment was studied by vacuum annealing (2 Pa) at temperatures (Ta) spanning the range 500 - 950 GC for a duration of one hour (ta„nCai). Prior work has shown that vacuum annealing is more effective in removing OH contaminants than annealing in air, as determined from increased Er+3 emission lifetime [13].

Film thickness and index of refraction were determined using a thin-film spectrophotometer. A model Fluorolog-3 spectrofluorometer (Horiba Jobin Yvon) was used to obtain room temperature photoluminescence (PL) emission from the Er+3 centers in the 1535 nm band. A Xe lamp was used for excitation with a double excitation monochromator to set a fixed excitation wavelength in the range of 515 to 530 nm. Spectral signal intensities were recorded in the range 1400 to 1600 nm using a LN2 cooled Hamamatsu InGaAs photodiode detector, preceded by a single emission monochromator, using 0.2 s integration time at each wavelength. Emission from Er+ in the 1535 nm band was maximized by tuning the excitation monochrometer to 521 - 523 nm, as obtained from a three-dimensional matrix of photoluminescence-excitation spectra. Photoluminescence spectra were normalized to the power output of the excitation source, monitored by a separate photodiode.

Annealing is generally required to density sol-gel films by removing water, organic compounds and hydroxyl residues [8,18]; most volatiles are driven off at T, ~ 500 °C [10]. Films may

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densify by -25%, observed in this work as shrinkage in thickness from -170 nm (air-dried) to -130 nm (Ta - 700 °C ). The Ta for producing the strongest PL depends on process method as well as Er concentration (e.g. Ta ~ 900 °C are often used for sol-gel films [6,7,21]). Decreased emission at high annealing temperatures can arise from a number of causes: Segregation by precipitation of Er or phase separation by clustering, and quenching by Er-Er interactions at high optically active Er+3 concentrations, as well as residual OH contamination [22].

Results: Erbium Emission

Photoluminescence spectra for samples annealed at various Ta are presented in Figure 1 (normalized logarithmic intensity scale). Peaks near 1160 nm originate from the silicon substrate (see section 3), while emission peaks near 1535 nm arise from Er+3 in the sol-gel film. Erbium emission from the as-deposited air-dried sample is very weak and hardly noticed when compared to annealed samples. This is to be expected, since low temperature baking leaves residues from the sol and an abundance of water and hydroxyls in a low-density porous near-glassy network. As such it strongly competes with the radiative ones, leading to the low PL signal observed.

The PL signal from the Er+3 appreciably improves as a function of increasing Ta, particularly for Ta > 700 °C, until it reaches a maximum at Ta « 850 °C. Samples annealed at higher Ta exhibit a decreasing trend in their PL (confirmed out to 1050 °C for additionally prepared samples annealed in either air or vacuum). Maximum PL intensity occurs at emission wavelengths in the vicinity of 1533 to 1537 nm and spectral full width at half maximum is in the range 51 to 58 nm, both varying somewhat with Ta. To illustrate the trend, peak PL intensities (normalized arbitrary units; integrated spectra are similar) are plotted against Ta in Figure 2 as filled circles (uncertainties for anneals above 600 °C are approximately 12% full scale). The evident non-monotonic dependence on Ta is indicative of competing thermal reactions that activate optical emission from the intra 4f Er+3 band as Ta is increased towards 850 °C, and de-activate it at higher Ta.

In [3], a model involving two activation energies was introduced to provide a quantitative analysis

is expected that non-radiative recombination

Figure 1. Photoluminescence intensity spectra from silicon coated with erbium-doped sol-gel silica films (anneal temperatures T„ in legend).

Figure 2. Maximum Er+Î photoluminescence intensities (-1535 nm) from silicon coated with erbium-doped sol-gel silica films vs. anneal temperature (data: points; theory: curve).

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of the dependence of the photoluminescence from Er on Ta; An energy Ea is defined as the activation energy of the favorable process (forming optically active Er+ ) and an energy Ed is defined as the activation energy of the unfavorable process (quenching optically active Er+3). Fitting the model to the data yields the results Ea = 2.3 ± 0.6 eV and Ed = 2.6 ± 0.7 eV, and the theoretical function shown in Figure 2; error bars in the activation energies arise from parameter correlation and uncertainty in the data. The magnitude of the activation energies is consistent with the idea that OH removal is the mechanism by which the otherwise optically active Er+3

yields the PL signal. Interpretations of these results are discussed below.

Influence of Hvdroxvl on Er13- Emission

Presence of OH is extremely effective in quenching excited Er+3 ions [17]. At high OH concentrations, direct resonant energy transfer from the excited ion to OH serves as an extremely effective trap. At low OH concentrations, fast transfer of energy from ion to ion via the Förster mechanism (cooperative energy transfer or CET) may allow diffusion of the excitation to an OH impurity, where it becomes dissipated by multi-phonon assisted decay to the oxide host [23]. The latter process is a type of concentration quenching, since the rare-earth concentration is the dominant factor in the CET coefficient [22], In the present case, both mechanisms are expected to apply, since the Er concentration is quite high.

Weak PL emission for the low-temperature anneals (Ta « 850 CC) indicates quenching of excited Er+3 ions by direct Er-»OH interactions. Increase in PL intensity in the region T a < 850 °C is therefore consistent with removing OH from proximity to Er+3 ions. Residual OH is mostly surface bonded within pores in the form of silanol groups, as shown by infrared absorption near 3670 cm"' observed in bulk silica gels (~ 900 °C anneals) [24,25]. An identified mechanism for OH removal is surface desorption of H2O, which is released upon reaction of surface silanol to form siloxane bridge structures, corresponding to an enthalpy AH = 2.16 eV [26]. (Bulk studies indicate that OH diffuses very rapidly at 850 °C [27].) Since E„ « AH within error, one may associate creation of optically active Er+3 with the process of OH removal by desorption.

Since the observed diminution of PL intensity at higher annealing temperatures corresponds to a thermal activation energy Ed that is also on the order ofthat for OH processes, it appears that OH removal is involved in the PL deactivation as well. Reduced concentrations of residual OH obtained for high temperature annealing may allow the diffusive Er—»Er—»OH processes to become effective for quenching PL emission. Moreover, high Er concentration in itself (e.g. upon virtually eliminating OH and densifying as Er-Er distances decrease) can lead to concentration quenching by cross relaxation or cooperative upconversion processes (CUP) involving Er-Er dipolar interactions, a form of CET, that typically converts one out of two units of excitation energy into heat; this CUP mechanism has been reported for optical fibers with high Er concentrations [28]. Both of these mechanisms can account for the decrease in PL emission for annealing in the region 850 °C « Ta < 1050 °C.

Diffusivity of Er at low concentrations in various deposited silica films was recently measured in the temperature range of 1000 - 1100 °C [29], from which an activation energy EEr = 5.3 eV was derived, due mainly to strong Er-O bonds, which are also responsible for stability of Er-glass structures; extrapolation to 950 °C yields diffusivity DE, « 6xl0"17 cmV and diffusion length (4DEr tançai f' ~ 10 nm. Erbium diffusion therefore appears to be sufficient for Er clustering at

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950 °C, owing to the high Er concentration in the sol-gel films (mean Er-Er distance ~ 1 nm); Er clustering generally kills the optical activity of the involved Er ions. Although Er segregation may participate along with the other mechanisms leading to overall PL quenching for Ta post 850°C, Er is considered to be a primary network component at comparable concentrations [15] favoring stable Si-O-Er glassy network.

As in fiber amplifiers doped with Er at high concentration, one expects that Er-Er interaction to saturate the PL but not to shrink it. However, as sol-gel films are usually not completely densified and Er diffusion is negligible for anneals near 850 °C, CET is minimized for Ta ~ 850 °C while OH out-diffusion is very welcomed. Higher annealing temperature permits the Er-Er interaction to occur at much higher rate due to faster densification process where PL deterioration can be spurred by anyone or more of the described processes.

These results are quite different from the concentration quenching observed at low Er concentrations. Existence of a range in annealing temperatures where concentration quenching may appear suppressed is a particular advantage obtained by using a high Er concentration. One does not expect thermal annealing behavior to materially depend on Er concentration, since the optimal anneal temperature depends mainly on the OH concentration that dominates the process. The optimum Ta at 6 at. % Er is somewhat lower than the previously reported value for sol-gel films with less than 1 at. % Er [6,7,21]; this is similar to the behavior of Er-doped silicon-rich oxides, where optimal Ta is lowered by about 100 °C at Er concentrations of 3 - 6 at. % [30].

3. Enhanced Optical Emission from Silicon

The PL spectra of Figure 1 shows emission at 1.067-eV (1162-nm), which is just below the Si band gap (energy units used herein by convention). Analysis presented below shows that this room temperature emission from silicon is enhanced by the Er-doped sol-gel coatings and is correlated with inhomogeneous film stresses. Obtaining high efficiency in light emission from silicon at room temperature has traditionally entailed selecting structures and materials to circumvent inherent disadvantages of the indirect Si band gap (e.g. typical 10"4 quantum efficiency at 300 K. [31], see also [32]) [33]. Thus, this sol-gel process method has inherent advantages as a simplified process.

Results: Enhanced Silicon Emission

Photolummescence spectra are shown in Figure 3 for two samples that were annealed at Ta of 700 °C and 850 °C (solid and broken curves, respectively) and exhibit major peaks at 0.807 and 1.067 eV. The 1.067-eV peak is strongest for annealing at 700 °C, where it is enhanced by a factor of 50 when compared to unannealed, air-dried films; as verified below, emission at 1.067 eV is associated with the Si substrate.

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The 0.807-eV peak is strongest for annealing at 850 °C (left dot-dash curve reduced by factor of 6), and is associated with emission from Er in the silica film, as discussed in section 2. Weaker emissions in the region 1.25 - 1.35 eV are attributed to 4ln/2 -> 4115/2 transitions in Er+3. Annealing at Ta = 850 °C corresponds to maximum 0.807-eV Er+3 emission. On comparing the two emission signals, the peak emission from the Si (Ta = 700 °C) is about 16% of the peak emission from Er+3 (Ta =850 °C); from areas under the respective spectra, the integrated signal from the Si is 50 % ofthat from Er+3.

The dotted curve in Figure 3 is a theoretical one-phonon (i.e. dominant phonon) model for emission of photons at energy E by phonon-assisted recombination of free electrons and holes in silicon, represented by the expression, PL(E) oc (E-E0)2 exp[-(E-Eo)/kßT (uncorrelated pair model, see e.g. [34]); it is calculated with threshold energy E0 = 1.020 eV and thermal energy kBT = 0.023 eV. Owing to this function, the PL spectra generally exhibit exponential tails for E>E0. Since the dominant phonons are the momentum-conserving transverse optical phonons of energy Eph = 0.0578 eV, the effective band gap, neglecting possible corrections for exciton binding or trapping, is EG = E0 + Eph = 1.078 eV, which is about 42 meV below the intrinsic Si band gap (1.12 eV). This is an approximate examination of the data as the effect of weaker phonon components, e.g. features appearing below 1.03 eV, were not included; the reasonable overlap for much of the PL signal peaking at 1.067 eV indicates that this emission originates indeed from the silicon substrate and not the film. X-ray diffraction data show no evidence of polycrystalline Si (e.g. absence of Si nano crystals within the deposited film [33]).

Annealing behavior for emission from Er+3 in the sol-gel film and from the Si substrate obey distinctly different dependences on Ta (compare Figure 2 and Figure 3, inset). This shows that emission from the Si does not depend on that from the Er+3 ions in the silica film, from which we conclude that emissions at these two wavelengths arise from independent mechanisms.

Influence of Sol-Gel Coating on Si Emission

The observed annealing behavior points to enhancement of emission from the silicon arising from a reversible (owing to non-monotonic dependence on Ta) interaction with the sol-gel coating. Stresses in sol-gel films, caused by shrinkage and porosity, are known to be inhomogeneous as indicated by IR absorption signatures (frequency shifted TO3 modes at 1060 - 1080 cm"1 [35]), and tend to be greatest for Ta ~ 700 °C [36], which coincidentally is where the strongest Si PL from our samples is observed. Given that Si PL is comparatively subdued for Ta < 600 °C as well as for Ta > 900 °C, one may conclude that enhanced PL at Ta ~ 700 °C is to be associated with non-uniformities in stresses in the films.

Figure 3. Photoluminescence intensity spectra for two silicon coated with erbium-doped sol-gel silica films, annealed at 700 °C (Si: solid curve) and 850 °C (Er+3: dashed; chain-dashed, scaled by 6). Dotted curve denote 1-phonon theory. Insert: peak PL from Si (-1.067 eV) vs. anneal temperature.

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Inhomogeneous film stresses, acting in conceit with native interfacial roughness, is therefore expected to create local strain-induced band bending in the silicon and increase the effective cross section for radiative free-carrier recombination. Applying this interpretation, enhanced emission is expected to be produced in only a thin layer at the Si surface (on the order of die film thickness, -0.1 um, or less); since the excitation radiation penetrates ~1 urn, quantum efficiency (QE) could actually be increased a thousand-fold (i.e. on the order often times larger than the observed enhancement), suggesting QE >1% (assuming QE < 0.01% as reported for CZ silicon [31,32]).

4. Conclusions

A process for producing optically active SiC>2:Er thin films on Si substrates using low cost sol-gel techniques that utilize Er2Û3 has been presented and analyzed. Photoluminescence is enhanced strongly as a function of annealing temperature, reaching optimum for annealing temperatures in a 50-°C range near 850 °C. External emission from Er-O structures is shielded by OH for annealing temperatures Ta < 850 °C and by the combination of Er concentration quenching and Er-Er-OH energy diffusion for Ta > 850 °C. The results indicate that minima in quenching by Er-Er cooperative energy transfer and by Er-OH interactions can be associated with Ta * 850 °C.

In addition, near-band gap (1.067 eV) photoluminescence at room temperature has been observed in CZ Si wafer material with the erbium-doped sol-gel silica coating. Emission from the Si, strongest for Ta « 700 °C, correlates with inhomogeneous stresses in sol-gel films (owing to strong 25% shrinkage), as indicated indirectly from IR absorption. Based on the independent annealing behavior of the two PL signals and comparison with prior studies of sol-gel film properties, it appears that the Er is not directly involved in the Si emission. The spin-coating process presented herein is notable for enhancing light emission from bulk-type Si (estimated QE ~1%) in the absence of any patterning or p-n junctions.

Acknowledgements

Partial support by the University of Jordan, the New Jersey Institute of Technology, and the U.S. National Renewable Energy Laboratory are gratefully acknowledged.

References

[1] M. Ferrari 2005 Handbook of Sol-gel Science and Technology, vol II, Processing, Characterization, and Applications, ed S. Sakka (CITY: Klewer Academic Publishers) pp. 359-389

[2] E. Desurvire, 1994 Erbium Doped Fiber Amplifiers: Principles and Applications, (New York, NY: Wiley) Chap 4

[3] S. Abedrabbo, B. Lahlouh and A. T. Fiory, J. Phys.D: Applied Phys. 44 315401 (2011) [4] S. Abedrabbo, B. Lahlouh, S. Shet and AT. Fiory, Scripta Materialia 65 767 (2011) [5] Q. Wang, N. K.. Dutta and R. Ahrens, J. Appl. Phys. 95 (2004) 4025-4028 [6] G.C. Righini, et al, Optical Quantum Electron. 34(2002) 1151-1166 [7] X.Orignac, D. Barbier, X. M. Du, R. M. Almeida, O. McCarthy and E. Yeatman, Optical Mater.

12(1999)1-18

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[8] Q. Xiang, Y. Zhou, B. S. Ooi, Y. L. Lam, Y. C. Chan and C. H. Kam, Thin Solid Films 370 (2000)243-247

[9] Y. Y. Hui, P-H. Shin, K-J. Sun and C-F. Lin, Thin Solid Films 515 (2007) 6754-6757 [10] C. K. Ryu, H. Choi and K. Kim, Appl. Phys. Lett. 66 (1995) 2496-2498 [11] X. Orignac, D. Barbier, X. M. Du and R. M. Almeida, Appl. Phys. Lett. 69 (1996) 895-897 [12] L. H. Slooff, M. J. A. de Dood, A. van Blaaderen and A. Polman, Appl. Phys. Lett. 76 (2000)

3682-3684 [13] L. H. Slooff, M. J. A. de Dood, A. van Blaaderen and A. Polman, J. Non-Cryst. Solids 296

(2001)158-164 [14] K. P. Huber and G. Herzberg, Molecular spectra and molecular structure IV. Constants

Diatomic molecules, (New York, NY: Van Nostrand Reinhold) (2001) 212 [15] J. T. Kohli and J. E. Shelby, Phys Chem. Glasses 32 (1991) 67-71 [16] J. E. Shelby and J.T. Kohli, J. Am. Ceram. Soc. 73 (1990) 39-42 [17] A. J. Bruce, W.A. Reed, A. E. Neeves, L.R. Copeland and W. H. Grodkiewicz, Proc. Optical

Waveguide Materials, Mat. Res. Soc, (Pittsburgh, PA: MRS) (1992) pp. 157-161 [18] S. Pal, A. Mandai, G. De, E. Trave, V. Bello, G. Mattei, P. Mazzoldi and C. Sada, J. Appl.

Phys. 108(2010)113116 [19] S. Wolf S, Silicon Processing for the VLSI Era, Process Technology, 2nd Ed., vol 2, (Sunset

Beach, California.: Lattice Press) (2000) Chap 16 [20] C. J. Brinker and G.W. Scherer, J. Non-Crystalline Solids 70 (1985) 301-322 [21] A. C. Marques, R.M. Almeida, A. Chiasera and M. Ferrari, J. Non-Cryst. Solids 322 (2003)

272-277 [22] A. Polman , Physica B 300 (2001) 78-90 [23] V. P. Gapontsev, S. M. Matitsin, A. A. Isineev and V.B. Kravchenko, Optics Laser Technol. 14

(1982)189-196 [24] S. Kondo, F. Fujiwara and M. Muroya, J. Colloid Interface Sei. 55 (1976) 421-430 [25] J. H., Jr. Anderson and K. A. Wickersheim, Surface Sei. 2 (1964) 25-60 [26] C. J. Brinker, G.W. Scherer and EP Roth, J. Non-Crystalline Solids 72 (1985) 345-368 [27] K. M. Davis and M. Tomozawa, J Non-Cryst Solids 185 (1995) 203-220 [28] R. Wyatt, Fiber Laser Sources and Amplifiers: Proc. SPIE vol 1171 ed M J F Digonnet

(Billingham, Washington: SPIE) (1989) pp 55-64 [29] Y-W. Lu, B. Julsgaard, MC. Petersen, R. V. S. Jensen, T.G. Pedersen, K. Pedersen and A. N.

Larsen, Appl. Phys. Lett. 97141903 (2010) [30] G. W. Adeola, H. Rinnert, P. Miska and M. Vergnat, J. Appl. Phys. 102 053515 (2007) [31] D. J. Lockwood, in: DJ. Lockwood (Ed.), Light Emission in Silicon: From Physics to Devices,

Semiconductors and Semimetals Series, vol. 49, Academic Press, Chestnut Hill MA, 1998, Chap. 1, pp 1-35.

[32] O. King, D. G. Hall, Phys. Rev. B 50 (1994) 10661-10665 [33] L. Pavesi and D. Lockwood, Silicon Photonics, Top. Appl. Phys. 94, Springer, Berlin, 2004. [34] G. Weiser, S. Kazitsyna-Baranovski, R. Stangl, J. Mater. Sei.: Mater. Electron. 18 (2007) S93-

S96. [35] P. Innocenzi, J. Non-Cryst. Solids 316 (2003) 309-319. [36] T.M. Panrill, J. Mater. Res. 9 (1994) 723-730.

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Influence of Annealing on the Martensitic Transformation and Magnetocaloric Effect in Ni49Mn39Snn Ribbons

Dianzhen Wu1'2, Sichuang Xue1'2, Wu Wang1'2, Qijie Zhai2, Hongxing Zheng1'2

1 Laboratory for Microstructures, Shanghai University, Shanghai 200444, China 2 Shanghai Key Laboratory of Modern Metallurgy & Materials Processing, Shanghai University,

Shanghai 200072, China

Keywords: Ni-Mn-Sn, Martensitic Transformation, Magnetocaloric Effect, Melt Spinning, Annealing

Abstract

Magnetocaloric effect (MCE) associated with martensitic transformation (MT) can be induced in Ni49Mn39Sni2 alloy. The influence of annealing on the martensitic transition and magnetocaloric effect in metamagnetic polycrystalline Ni49Mn39Sni2 ribbons produced using melt spinning technique was investigated systemically based on the results of vibrating specimen magnetometer (VSM) measurements. Ribbons were annealed at 1273 K for 4h, followed by water quenching or furnace cooling to room temperature. Magnetic-field induced reverse MT from paramagnetic martensite phase to ferromagnetic austenite phase together with large magnetization change is evident in thermal-magnetic curves, thus entropy change appears. The magnetic entropy change enhances strikingly for all annealed ribbons in comparison to melt-spun ribbons. The magnetic refrigerant parameters for water quenched ribbon exhibits optimum values; it has extensive potential applications in practical refrigerant cycles.

Introduction

Magnetic refrigeration, which takes magnetic materials as working substance, is highly efficient and protective for the environment in comparison with vapor refrigeration. It utilizes cycle of magnetization and demagnetization. The key for large scale applications lies in cheap materials with giant magnetocaloric effect (MCE) under low magnetic field. Among the families of magnetic refrigerant materials, ferromagnetic shape memory alloy, Ni-Mn-Sn, has been receiving increasing attention due to large MCE and low cost [1-3]. However, most recent studies focus on bulk alloys with indispensable long-duration high-temperature anneals [4], Melt-spinning technique has been proven to be an effective method to synthesize single-phase materials with homogeneous chemical composition, avoiding or reducing the time of annealing significantly [5, 6].

Experimental procedures

An as-cast Ni49Mn39Sni2 (at.%) ingot was repeatedly arc melted from Ni, Mn, Sn with purities of 99.99 wt.% under argon gas atmosphere. Additional 5 wt.% Mn was added to compensate for evaporation losses. The arc-melted sample was induction melted in a quartz tube with a rectangle nozzle (6 mm x 0.5 mm) and then ejected onto a copper wheel rotating at a linear speed of 10 m/s. The resulting as-spun ribbons were sealed in quartz tubes filled with argon gas and annealed at 1273 K for 4 h followed by water quenching or furnace cooling, hereinafter referred to as "as-spun ribbon", "ribbon A" and "ribbon B", respectively.

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The crystal structure was identified by X-ray diffraction (XRD DLMAX-2550) using CuKa radiation. Magnetic measurements were performed using a vibrating sample magnetometer (VSM LAKESHORE 7404) under a magnetic field up to 1.8T.

Results and discussion

X-ray diffraction patterns for all three ribbons obtained at room temperature are shown in Fig. 1. It is obvious that all the samples have shown a single bcc cubic L2i structure, with lattice constant presented in Table 1. What is more, no secondary phases are detected in the XRD patterns.

Figure 1. XRD patterns for as-spun ribbon (a), ribbon A quenched from 1273 K (b), ribbon B furnace cooled from 1273 K (c), respectively. All ribbons in the present study are austenitic at room temperature.

Fig. 2 presents the thermomagnetic M (T) curves for Ni49Mn39Sni2 ribbons under a magnetic field of IKOe. It can be observed that all three samples show similar tendency. Decreases in magnetization are first observed with heating; afterwards a sudden jump appears due to the reverse MT from paramagnetic martensite to ferromagnetic austenite. As temperature increases further, the austenitic phase transforms to paramagnetic state at the Curie temperature TC

A. The martensitic start and finish temperature (As and Af) and TcA, defined as the temperature where dM/dT obtain maximums, are denoted in the figure (just shown in ribbon A for clarity). All the characteristic temperatures are summarized in Table 1. We can find that structural transition temperatures increase remarkably after annealing, while TcA keeps almost constant. Additionally, for the annealed ribbons, magnetization of austenitic phase enhances, while martensitic phase weakens, thus leading to relatively large magnetization discrepancy (AM) between the two phases, especially ribbon A. The strengthened AM suggests preferred magnetic entropy change (AS).

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Figure 2. Temperature dependence of magnetization for all ribbons in a field of 1 kOe.

Fig. 3 shows the isothermals measured in the vicinity of MT and second-order magnetic transition, by means of increasing and decreasing the field of 1.8T. For a better clarity, only some specified plots around MT are selected in Fig. 3(a), (c) and (e). Metamagnetic behavior, denoted by arrows, owing to the magnetic field induced structural transition can be observed for all samples. Consequently, noticeable hysteresis loss, shown as the stripped area in Fig. 3, turns to be reasonable, while for the austenite, no metamagnetic transition and hysteresis appear.

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Figure 3. Magnetization isothermals in the vicinity of first-order MT [(a), (c) and (e)] and second-order magnetic transition [(b), (d) and (f)] for all ribbons. The stripped areas represent the hysteresis loss due to magnetic field-induced reverse MT.

Hysteresis losses in the 1.8T field as a function of temperature are the calculated areas between magnetization and demagnetization branches of isothermal curves, plotted in Fig. 4. They are negative factors which weaken the refrigerant efficiency significantly near the first-order MT, and consequently, should be computed in a proper way. The estimated average hysteresis losses are calculated by integrating the area under the curves. It is obvious that the magnetic hysteresis of annealed ribbons turn out to be extremely larger in comparison to as-spun ribbon. For ribbons A and B, the one cooled slowly has a relatively small hysteresis.

Figure 4. Hysteresis loss as a function of temperature around MT in the magnetic field of 1,8T for all ribbons.

Magnetic entropy change (AS) and refrigerant capacity (RC) are two important parameters to estimate the extent of MCE at phase transition. AS as a function of temperature for all samples are calculated employing the Maxwell relationship and is shown in Fig. 5. The Maxwell relationship can be approximately written as in equation (2) [7], The value of AS for the average temperature refers to the integrated area between two adjacent magnetization isothermals from 0 to 1.8T. It clearly shows that AS peaks around MT and second-order magnetic transition exhibits different features. The peaks with high values at the structural transition are sharp over narrow temperature ranges. For the magnetic transition, AS changes its sign and drops significantly over a broader temperature range, due to the character of second-order transition, as shown in Fig. 5(b). However, as-spun ribbon is an exception, as the AS of MT is close to magnetic transition. Moreover, entropy change, AS for the MT, strengthens heavily after heat treatment with ribbon A being especially high. It should be noted that heat treatment has little effect on AS around

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magnetic transition. For the structural transition of ribbon A, the maximum value of AS (ASM) is 13.17J/(kg-K), which is higher than that of bulk NisoMnsySnn [2] and Niso3Mn35 5Sni4 4 [3] alloys.

ASlt(T,H) = MojjdM/dr)HdH (1)

ASM = ̂ F [I" M{T+AT'H*>dH ~ §t(T,H)<m\ (2)

Figure 5. Temperature dependence of magnetic entropy change in the vicinity of (a) MT and (b) second-order magnetic transition.

In addition to entropy change, the refrigerant capacity RC is another important parameter, meaning the amount of heat that can be transferred in a real refrigerant cycle. RC refers to the integrated area under AS peaks, and is defined as [8]:

RC = £[ASu{T))mdr (3)

where, Ti and T2 correspond to the temperature at the full width half maximum peak value of AS. Additionally, the temperature range between Ti and T2 is the real working range of the material. As for the refrigerant materials used near room temperature, Erikson cycles which require wide working range are needed. It is worth noting that larger RC value is available for the magnetic transition, in spite of lower AS value. For the first-order MT, RC is smaller, though the maximum AS peaks are higher. It is reasonable considering the broader AS temperature dependence for the magnetic transition and sharp peaks of MT. In other words, the working temperature range is a decisive factor for RC value. As mentioned above, hysteresis loss is an unwanted character that hinders RC. It needs to be subtracted so that we can obtain the effective refrigerant capacity RCzn around MT. It must be pointed out that magnetic transition is more preferable than MT as a result of no hysteresis, larger RC values and broader working temperature range. Especially for ribbon A, RC achieved around magnetic transition gets to be almost twice that of MT. What is more, it has been reported that both the magnetic transition and MT can be applied in a refrigerant cycle. Consequently, RCM^ which is the sum of Rda and RC at 7^ turns to be the total refrigerant capacity of a real cycle. The relevant magnetocaloric parameters are also listed in Table 1. It can be suggested that the total effective RC obtained on an entire refrigerant cycle for all samples, especially ribbon A, could reach 69.97J/kg. This is comparable to bulk alloys of Ni-Mn-Ga [9], Ni-Mn-In [7] and melt-spun ribbons of Ni50.3Mn35.5Sn 14.4 [3],

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It has been found that heat treatment influences significantly the magnetocalonc performance of ribbons. As we know, the atomic site in as-spun ribbons tends to be unstable due to direct solidification from melt to the solid state with extremely large cooling rate. There is little time to take up the equilibrium positions corresponding to lower energy. After heat treatment, atoms jump to more balanced positions with lower energy, and then atomic sites may be adjusted slightly, allowing the existence of modification of Mn-Mn distance, which can be demonstrated by the change in lattice parameters for different ribbons as mentioned above. Magnetic interactions are closely related to Mn-Mn distance [10], and consequently for change of magnetic properties, such as transition temperatures and magnetocaloric performances.

•Sample« Parameter ÂtKÏ À,IK\ c AC at 7^ RC at 7 M hvstirfiii RC°" RC'M

Samples Parameter A. (K) * ( K ) ( R ) ( J / k g ) ( J / k g ) by*«» ( J / k g ) ( J / k g )

as-spun

ribbon 5,987 242 270 308 15.20 14.80 1.79 13.01 28.21

Ribbon A 5.917 262 280 310 49.06 33.75 12.84 20.91 69.97

Ribbon B

Table 1. Lattice parameter, transition temperatures, and magnetocaloric parameter around the MT and 7? for a field change of 18 kOe.

Conclusions

We have studied the effects of annealing on the martensitic transition and magnetocaloric properties of Ni49Mn39Sni2 ribbons. The influence of annealing can be ascribed to the change in Mn-Mn distance. Ribbon A has the optimal MCE in comparison to other two samples. The maximum value of entropy change around MT and total effective RC for a magnetic field change of 0-1.8T reach 13.17J/(kg-K) and 69.97J/kg, which are comparable to bulk alloys. As refrigerant materials, Ni49Mn39Sni2 ribbons have several advantages. Additionally, transition temperature range approaches room temperature, and is indispensable for practical refrigerant applications. Furthermore, the low cost of raw materials and convenient fabricating methods make ribbon A a promising potential magnetic refrigerant material.

Acknowledgements

All the authors gratefully acknowledge the support from the Shanghai Pujiang Program (10PJ1404200), the Scientific Research Foundation for the Returned Overseas Chinese Scholars, State Education Ministry and the Innovative Foundation of Shanghai University of China.

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References

1. Thorsten Krenke, et al., "Martensitic transitions and the nature of ferromagnetism in the austenitic and martensitic states of Ni -Mn-Sn alloys", Physical Review B, 72(2005), 014412.

2. Thorsten Krenke, et al., "Inverse magnetocaloric effect in ferromagnetic Ni-Mn-Sn alloys", Nature Materials, 4(2005), 450-454.

3. B. Hemando, et al., "Magnetocaloric effect in melt spun Niso3Mn35sSni44 ribbons", Applied Physics Letters, 92(2008), 132507.

4. D.L. Schlägel, et al., "Influence of solidification microstructure on the magnetic properties of Ni-Mn-Sn Heusler alloys", Journal of Alloys and Compounds, 463 (2008), 38—46.

5. I Babita, et al., "First order structural transformation and inverse magnetocaloric effect in melt-spun Ni-Mn-Sn ribbons", Journal of Physics D: Applied Physics, 43(2010), 205002.

6. H. C. Xuan, et al., "Effect of annealing on the martensitic transformation and magnetocaloric effect in NJ44.iMa^Sniu ribbons", Applied Physics Letters, 92(2008), 242506.

7. Chao Jing, et al., "Exchange bias behavior and inverse magnetocaloric effect in NisoM^sInis Heusler alloy", Journal of Alloys and Compounds, 475 (2009), 1-4.

8. K.A. Gschneidner, et al., "Recent Developments in Magnetic Refrigeration", Materials Science Forum, 315-317(1999), 69-76.

9. Shane Stadler, et al., "Magnetocaloric properties of Ni2Mni-xCuxGa", Applied Physics Letters, 88(2006), 192511.

10. Yu SY, et al., "Realization of magnetic field-induced reversible martensitic transformation in NiCoMnGa alloys", Applied Physics Letters, 91(2007), 102507.

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Metal Diaphragm Based Magnetic Field Sensor

Asahel Banobre, Ivan Padron, Anthony T. Fiory, Nuggehalli M. Ravindra

Department of Physics, New Jersey Institute of Technology, Newark NJ 07102, USA

Keywords: Magnetic Field Sensor, Metal Diaphragm, Diaphragm Based Sensor

Abstract

A diaphragm based magnetic field sensor is designed and fabricated. A thin metal diaphragm is used as a sensing element and it responds to an external magnetic field. Deformation of the diaphragm due to the pressure produced from the interaction between the magnetic field and the metal surface is detected electronically through a Wheatstone bridge placed over the diaphragm. Analytical and experimental analyses are used to study the precision and accuracy of devices for sensing magnetic field flux.

Introduction

The measurement of magnetic fields has been indispensable for many technical areas such as defense systems, geophysics, industry and engineering applications [1]. Currently, not one factory, automobile, or computer could operate as efficiently without magnetic sensors; magnetic sensors have been used for medical and biological applications: magnetic tags can be attached to detect the presence of specific molecules, SQUIDS have been used to measure extremely weak magnetic fields generated by the brain [2] and in other medical tests [3].

There are several technologies to detect the presence, strength, or direction of magnetic fields not only from the Earth, but also from permanent magnets, from magnetized soft magnets, and from the magnetic fields associated with current. These sensors are used as proximity sensors, speed and distance measuring devices, navigation compasses, and current sensors [4], Fluxgate, Hall effect, and magnetoresistivity are some of the techniques used; each of them have a different approach to using magnetic sensors. Each type of technology focuses on a particular area for detection, a measurement to be detected and a way of recording changes. There is not much need to improve the sensitivity independent of size, power, and cost. Instead, for each application, one needs to make a tradeoff between sensitivity, size, power, and cost. A major need is to reduce the cost of the signal processing electronics since, in many cases, the signal processing electronics is much more expensive than the sensor element. [1]

MEMS magnetic field sensor has been developed since the 1990s to reduce the size of many types of sensors as well as to increase their ability to be mass produced. Recently, the magnetic field sensors using MEMS technology have been reported in several studies. [5-8]

The metal diaphragm based sensor as part of die MEMS technology will allow the integration of magnetic field sensors with electronic components, which offers important advantages such as small size, light weight, minimum power consumption, low cost, better sensitivity and high

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resolution, improving the compromise between the sensitivity, size, power and cost, as well as enhance its utility in applications involving corrosive media and high temperatures.

Sensor Design and Fabrication

We propose to use a thin metal diaphragm with a Wheatstone bridge configuration as sensing element to respond to an external magnetic field. When the sensor is subjected to an external magnetic field, the magnetic force bends the metal diaphragm, producing stress and resistance change in the Wheatstone bridge and the output voltage of the sensor.

A schematic illustration of the metal diaphragm magnetic sensor proposed in this studied is shown in Figure 1. The sensing elements comprise a metal circular diaphragm (Figure 2 (a)) and strain gauges bonded to the diaphragm in a Wheatstone bridge configuration. The interaction of an external magnetic field with the metal generates a magnetic force which bends the diaphragm. The piezoresistive strain gauges (p-type) are sensitive to the strains induced by diaphragm deflection. Changes in the resistance of the gauges are detected by a Wheatstone bridge circuit which yields an output signal.

Figure 1. Configuration of the metal diaphragm based magnetic sensor.

Conventional fabrication techniques [9, 10] were used to fabricate the circular diaphragm; it is made from a stainless steel 300 series sheet (Precision Brand Product, Inc.) which is milled until the desired thickness is attained. The strain gauges (commercially available) are directly bonded to the diaphragm and interconnected in a fully active Wheatstone bridge configuration, as shown in Figure. 2(b). In response to the induced strains due to the diaphragm deflection, a change in output voltage from the Wheatstone bridge scales with the external magnetic field intensity.

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Figure 2. Circular metal diaphragm with a Wheatstone bridge configuration.

Circular Diaphragm Theory

The mathematical analysis of the circular diaphragm will be based on a flat diaphragm with uniform thickness and composed of isotropic and homogeneous material, where the maximum deflection due to a normally applied pressure (due to external magnetic field) will not be more than 30% of the thickness. The deflection is due mostly to bending and the diaphragm will not be stressed beyond its elastic limit, and all loads are applied normally [11-14].

The maximum deflection, Wo. of a circular diaphragm with fixed edges loaded by pressure P occurs at the center and is:

3(1-M»)P.« ( 1 )

where a is the diaphragm's radius and h the thickness of the diaphragm; fi and E are Poisson's ratio and Young's modulus, respectively.

At the center, the maximum stresses c w correspond to:

3 Pa2

<W = ±g(l+fl)-£T (2)

Wheatstone bridge output

For small deflections, the stress can be taken as a linear function of the applied pressure differential across the two sides of the membrane. The output from the Wheatstone bridge circuit shown in Figure 1(b) is given by the expression [15-18]:

I R\ R4 \_AR Vout " V" \RT+T2 ~RT+FJ ~ 1^ (3)

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where, Vt, is the bridge voltage (input voltage) and R{ corresponds to each of the bridge piezoresistor. Since the fractional change in resistance is proportional to the stress:

AR GFa

where, Rg is the resistance of the un-deformed gauges, G/r the Gage Factor and £ = - the strain. Combining equation 2, 3 and 4, we obtain the expression for the linear relation between the output and the pressure:

/3GF(1 + n)a2\ V«*=V»{ g j )P + C &

where the offset C is introduced to take into account zero offsets in the Wheatstone bridge (imperfect resistor matching).

Magnetic Sensitivity

The magnetic force acting over the area A of a circular diaphragm due to an external magnetic field flux B can be approximated according to the following relationship between the magnetic force and the magnetic flux intensity B [5]:

F=^-B2A (6)

where, F is the magnetic force, fi0 is the permeability of free space, and B is the magnetic field intensity and A= a2 the area of circular diaphragm perpendicular to the magnetic field. Combining equations 5 and 6 with the definition of pressure as the force per unit area, we can relate the sensor output directly with the intensity of the external magnetic field as:

'3GF(l + /^)27^a2,

Vont = Vb | QEh2 "° ] ß 2 + C (7)

Experimental set-up

The experimental set-up was designed to accurately detect the intensity of an external magnetic field in perpendicular direction with respect to the flat metal diaphragm. A permanent magnet was used as the external magnetic field source. The use of a micrometer and a Hall Effect Gauss meter allows us to calibrate the magnitude of the magnetic field intensity with respect to the distance of the magnetic source and sensor.

Figure 3 shows the magnitude of the magnetic field intensity corresponding to the specific distance of the magnet (magnetic source) from the sensor.

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Figure 3. Magnetic field calibration.

Two stainless steel circular diaphragm based sensors were fabricated for this study; one of them (identified as sensor 1) was designed to operate in the pressure range of 0 to 25 psi and has overpressure tolerance of 50 psi, and the second (identified as sensor 2) was designed for a pressure range of 0 to 50 psi with an overpressure tolerance of 100 psi. Sensor 1 has a diaphragm thickness of 50.8 um and sensor 2 has a diaphragm thickness of 63.5 urn. The radius of both sensors was approximately 1.27 mm (1270 urn). The electronic reading was collected with 5 volt input on the Wheatstone bridge and the Gage Factor for the strain gages (p-type) used is 80.

As a result of the magnetic field calibration, our sensor was tested in the range of 1 to 10 millimeter due to the reading stability.

Sensor Output

Using equation 7, we can estimate the output voltage of the sensor 1 and sensor 2 due to the interaction with the external magnetic field; the Poisson's ratio and Young's modulus for stainless steel grade 300 series are 0.3 and 193 GPa, respectively (Table. I).

Table I. Sensors output (Calculated)

B(Gauss)

68 58

49 42 36 30

26 22

20 18

Sensor 1 Output (mV)

322 306

294 286 280 274

272

269

268 267

Sensor 2 Output (mV)

47 33

23 16 11 7

4 2

2 1

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Figures 4 shows the sensor output for both sensors; in each case, the experimental data differ from the calculated value (quadratic behavior). The deviation recorded in both figures comes mainly from the positioning error as well as the magnetic source (magnet) and Gauss meter used. Experiments with different magnetic sources and a more accurate Gauss meter are in progress to validate the theoretical calculation and improve the sensor performance.

Figure 4. Sensor output, experimental and calculated data.

Further work needs to be done to explore the dynamic performance and temperature response of the designed sensor.

Conclusions

We have proposed a stainless steel diaphragm based magnetic field sensor with a Wheatstone bridge configuration as sensing element to respond to an external magnetic field. As part of the MEMS based sensor, it offers important advantages such as small size, minimum power consumption, low cost, easy integration, better sensitivity and high resolution. Additional benefits include applications involving corrosive media and high temperatures.

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The preliminary data reported in the above study has confirmed the ability to introduce a metal diaphragm as an alternative technology for the design and fabrication of magnetic sensors. However, more work can be done to improve the sensor performance and implement new applications.

References

1. J. Lenz and A. S. Edelstein, "Magnetic Sensors and Their Applications," IEEE Sensor Journal, 6(3) (2006), 631-649.

2. S. Masahiro, T. Hiroaki, K. Kunio, and H. Yasuhiro, MEGvision Magnetoencephalograph System and Its Applications, 2004, vol. 38, pp. 23-27, Yokogawa Tech. Rep. (English Edition).

3. R.Wakai, "Current and future technologies for biomagnetism," in Proc.AIP Conf, 2004, Vol. 724, pp. 14-19.

4. M. Caruso, T. Bratland, C. Smith, and R. Schneider, "A new perspective on magnetic field sensing," Sensor Magazine, Dec (1998), 34-45.

5. D. Guangtao, C. Xiangdong, L. Quibin, L. Hui, and G. Huihui, "MEMS Magnetic Field Sensor Based on Silicon Bridge Structure," Journal of Semiconductors, 31(10) (2010), 10401 l-[l-6].

6. D. Ciudad D, C. Aroca, M. C. Sanchez, E. Lopez, P. Sanchez, "Modeling and fabrication of a MEMS magnetostatic magnetic sensor," Sens. Actuators A, Phys., 115, (2-3) (2004), 408-416.

7. V. Beroulle, Y. Bertrand, L. Latorre, P. Nouet, " Monolithic piezoresistive CMOS magnetic field sensors," Sens. Actuators A, Phys., 103 (1-2) (2003), 23-32.

8. H. H. Yang, N. V. Myung, J. Yee , D. -Y. PARK, B. -Y. Yoo, M. Schwartz, K. Nobe, and J. W. Judy, "Ferromagnetic micromechanical magnetometer," Sens. Actuators A, Phys., 97-98 (2002), 88-97.

9. S.-P. Chang, J.-B. Lee and M. G. Allen, "Robust Capacitive Pressure Sensor Array," Sensor and Actuators A, 101(2002), 231-238.

10. S.-P. Chang and M. G. Allen, "Capacitive pressure sensors with stainless steel diaphragm and substrate," J. Micromech. Microeng. 14 (2004) 612-618.

11. S. Timoshenko and S. Woinowsky-Krieger, Theory of Plates and Shells (McGraw-Hill, 1959).

12. Mario Di Giovanni, Flat and Corrugated Diaphragm Design Handbook (Marcel Dekker, Inc., 1982).

13. E. Ventsel and T. Krauthammer, Thin Plates and Shells. Theory, Analysis, and Applications (Marcel Dekker, Inc, 2001).

14. W. P. Eaton, F. Bitsie, J. H. Smith, D. W. Plummer "A New Analytical Solution for Diaphragm Deflection and its Application to a Surface-Micromachined Pressure Sensor," MSM 99,Technical Proceedings of the 1999 International Conference on Modeling and Simulation of Microsystems, (1999), 640 - 643.

15. S. M. Sze, Semiconductor Sensors (John Wiley & Sons, Inc., 1994). 16. U. Gowrishetty, K. Walsh, J. Aebersold and H. Millar "Douglas Jackson "Development

of Ultra-Miniaturized Piezoresistive Pressure Sensors for Biomédical Applications," University/Government/Industry Micro/Nano Symposium, 17th Biennial, (2008), 89 -92.

17. R. Singh, L. L. Ngo, H. S. Seng, F. N. C. Mok "A Silicon Piezoresistive Pressure Sensor," Proceedings. The First IEEE International Workshop on Electronic Design, Test and Applications, (2002), 181-184.

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18. R. Buchhold a, A. Nakladal, G. Gerlach, P. Neumann, "Design studies on piezoresistive humidity sensors," Sensors and Actuators B, 53 (1998), 1-7.

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Supplemental Proceedings: Volume 1 : Materials Processing and Interfaces TMS (The Minerals, Metals <£ Materials Society), 2012

Optical and Electronic Properties of AIN, GaN and InN: An Analysis

Chiranjivi Lamsal, Dongguo Chen and Nuggehalli M. Ravindra

Department of Physics, New Jersey Institute of Technology, University Heights, Newark, NJ 07102, USA

Optical properties, A1N, GaN and InN, Nitride Semiconductors

Abstract

Optical properties of III-V nitrides, A1N, GaN and InN, have been analyzed in the study. The analysis takes into account the available data of the optical properties in the literature. The related optical properties such as the energy dependent dielectric fonction, and temperature and pressure dependent band gap are analyzed. Penn-like models are deployed to interpret the peaks in the reflectivity spectra and are compared with available data in the literature. Structure-dependent dielectric functions are investigated from both experimental and theoretical perspectives.

Introduction

The III-V semiconductor crystals such as A1N, GaN, and InN are of enormous technological importance in semiconductor technology for many applications in optoelectronics [1]. Recently, III-V nitride semiconductors are used extensively to improve the performance of blue and green light emitting diodes (LEDs). These devices have wide applications in rapidly growing semiconductor technology such as optical-fiber communication networks, compact disk optical-storage technology, laser printers as well as in automobile industry. Each of the nitride materials has high ionicity, high bond strength and good thermal conductivity [2] and are candidates for high power and high temperature device applications. Wurtzite materials have direct energy gap ranging from 6.1 eV [3] for A1N, 3.5 eV [3] for GaN, and 0.8 eV [4] for InN. GaN and A1N are wide-energy gap materials; these materials have low dielectric constants and are important for optoelectronic devices [5]. Hardness, high melting point, high thermal conductivity and large bulk moduli are their characteristic mechanical properties which make them useful for protective coatings [6]. Owing to low energy gap and the capacity of nitride alloys to cover the whole solar spectrum, InN could be useful for fabricating future generation solar cells [7].

Temperature and pressure variation results in a change in the optical properties. The variation in band gap with parameters such as pressure or temperature is one of the characteristics of a semiconductor which is fundamental and very important for applications in device technology. Theoretically, it is known that most of the variation of the energy gap of semiconductors with temperature is due to the following two mechanisms: (i) The major contribution to the temperature dependence of the energy gap of semiconductors comes from a shift in the relative position of the valence and conduction bands because of a temperature dependent electron-lattice interaction, (ii) The temperature dependent dilatation of the lattice causes a shift in the relative position of the valence and conduction bands [8,9].

701

In order to interpret the frequency dependence of the dielectric functions ex and e2, Perm-like models are used in this study. Band structure of a material is related to its e1 —E spectrum and an intensity maximum obtained in the spectrum corresponds to a maximum number of optically induced electronic transitions [10]. Since the energy corresponding to the peak in the £j — E spectra is a "macroscopic" [11] band gap, there should be a relation between the energy gap and high-frequency dielectric constant £œ(« n2), where n is the refractive index. According to Penn model [11], the high-frequency dielectric function is given by:

«- = 1 + ^ 1 - ^ ) + ^ / « , « <D

or

i = ^+« (£"~%„p)2+;(V8zrP2]^ (2)

where, Ep is the Perm gap and EF is the Fermi energy given by [11]:

EF = 0.2947(fto)p)4 (3)

with valence-electron plasmon energy given by [13]:

hü)p = 28.8( erfP), W is the molecular weight and Neff is the number of valence electrons per atom calculated by using:

Neff =Ma + N(8 - b)

for a compound AMBN, where a(b) is number of valence electrons per atom of type A(B) and M(N) is the atomic fraction of element A(B).

Results and discussion

The optical properties of III-V nitrides i.e. AIN, GaN and InN are analysed in this study based on the results of the frequency dependence of optical constants n and k, in the literature [14-18]. The AIN thin film samples were deposited on ultra-high-purity fused quartz (glass) substrates [12] whereas the single-crystalline hexagonal GaN [16] and InN [17,18] films were grown on (0001) sapphire substrates by metal organic vapor phase epitaxy.

Fig. 1 shows the variations in the optical parameters n, k, e2 and R with energy for AIN, GaN and InN in the energy range 1.5-6.6eV, 1.2-9.4eV, and 2-10eV respectively. It is seen that these parameters are sensitive to the energy of the incident photons to a greater extent and show maxima at certain frequencies. The variation in the e2 - E spectra of AIN follows parabolic form of 0.08x2-0.45x+0.58 in the energy range 1.5-5.7 eV and the intensity increases rapidly to 4.4 at E=6.6eV with a kink at 6.25eV. The reflectivity spectrum follows a parabolic variation of the form 0.65x2-2.8x+16 until it reaches the kink at 6.15 eV with reflectivity 25% as shown in Fig. 2.

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In the case of GaN, e2 is practically zero until 3eV and increases abruptly to 1.8 at 3.6 eV. While there is an increase in the value of e2 until 4.6eV, the rate of increase is faster from 4.6eV to Ei=6.8eV until it reaches its maximum intensity of 9.5. A saddle point of intensity 6.5 is observed at energy E2=7.8eV and then decreases further and reaches a local minimum of 4.5 at energy 8.8 eV. The reflectivity on the other hand, increases with energy until it finds a shoulder with reflectivity of 22% at energy of 3.4eV but decreases to a value of 20% at 3.8eV. However, it increases and reaches its maximum value of 38% at an energy of 7eV. Finally, the reflectivity starts to decrease but finds a small "hill" at 8eV.

The variation of e2 in InN is similar to that of GaN with shoulder (intensity=3.1) and maximum peak value (intensity=5.9) occurring at 2.6eV and 5.2eV respectively; which are relatively low as compared to that of GaN. A local minimum of value 3.2 is observed at 7eV and two kinks separated by leV are observed at 7.8 and 8.8eV in the energy range of 7-10eV. On comparing with GaN, relatively similar energy dependence is observed in the reflectivity curve of InN with corresponding peaks shifted towards lower energy values. The highest peak with reflectivity 27% occurs at 5.2eV. This decrease in intensity and corresponding location may be attributed to low energy gap observed in InN. The reflectivity for photons with energy 7eV is minimum (17%) in the observed energy range of reflectivity for InN; two more kinks separated by nearly leV(0.8eV) are observed at higher energies with increasing reflectivity.

Energy (eV) Fig. 1 Variation of n, k, and R as a function of photon energy.

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Fig. 2 Variation of E2 and R as a function of photon energy for AIN

The extinction coefficient is sensitive to the energy of incident photons to a greater extent and show maxima at certain frequencies.

Energy gap determines the threshold for absorption of photons in semiconductors. The pressure dependence of the absorption edge is quantified by pressure coefficient of the gaps, defined as —^and is given for (wurtzite) GaN by the slope of a typical "sublinear" relation [19] (and Fig 3aP):

E0(p) = -0.0018 x p2 + 0.047 x p + EB(p = 0)

Where, p is in Gpa and E0 in eV. According to the experimental results of Perlin et. al. [19], energy gap Eg for GaN is 3.44 eV and pressure coefficient is 4.7 x 10"2 eV/GPa. They have concluded that the GaN main gap is direct contrary to the previous calculation that the gap is indirect [20].

Fig.3 Energy gap versus pressure for (a) GaN (b) AIN (c) InN

From the results of Akamura [21] (Fig. 3b), the pressure coefficient of direct band gap of AIN is found to be 49 meV/GPa. This value is close to the one calculated by first principle method which is 43 meV/GPa [19]. Van Camp et. al [22] has reported even smaller value which is 36.3 meV/Pa. Although AIN also has both wurtzite and zinc blende (ZB) forms, the studies of the

704

optical properties of ZB AIN has not been reported in the literature [23]. However, ab initio calculations [22] have shown that the AIN in ZB form is a semiconductor with an indirect band gap of 4.04 eV which is smaller as compared to its corresponding wurtzite counterpart of 6.1 eV [3].

From the result of Ibanez [24] [Fig.3c] the pressure coefficient of direct band gap of wurtzite InN is found to be 29 meV/GPa. Similarly the value of pressure coefficient has been reported to be 30meV/GPa by Li et. al [25].

The variation of energy gap with temperature in (wurtzite) GaN thin films (grown on C-AI2O3) is shown in Figure 4a [26].

Fig.4 Energy gap vs Temperature for GaN, AIN and InN

The triangle represents the results of Li et al. [26] which is fitted with the Varshni relation [27], which describes well the temperature dependence of both the direct and indirect band gap in a bulk semiconductor :

A T2

£(r) = £(0) -^- j -g

where, E(fl) is the energy gap at 0 K, and 0 and A are the Debye temperature and Varshni coefficients, respectively. The results of Li et al., shows that E(0) = (3.470 ± 0.003) eV, A = (5.9 ± 0.5) x 10"4 eV/K and 0 = 600 K [28]. ß - GaN has also direct gap at room temperature which is less up to 0.6 eV [29,30] than that of wurtzite GaN. Ramirez-Fbres et. al [31] show that for cubic GaN, A = 6.697 x 10~4 eV/K, E(0) = 3.302 eV where they have used same value of Debye Temperature i. e., 600 K.

The temperature dependence of the band gap in AIN is shown in Figure 4b [32]. Results of Nam et. al shows that E(0) = 6.066 eV, A = 2.59 meV/K and 0 = 2030 K. The Debye temperature of AIN is greater than that of GaN which is expected because of the larger band gap in AIN.

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The temperature dependence of band gap in InN is weaker. Wu et. aL [33] have reported that for wurzite sample, of InN films, the Varshni coefficient and Debye temperature has been estimated to be QAlmeV/K and Ö = 454 K respectively, and f(0) is found to be 0.69 eV.

In Figure 4c, the solid line is the fit to the Varshni relation compared with the result of Wu et. aL

Table 1: Optical parameters for BN, AIN, GaN, and InN. Parameter E«(0)eV A (meV/K)

em dEa -zS- (meV/GPa) dp Debye Temp (K)

Melting Temp (K)

Wurtzite

ZB

Wurtzite

ZB

BN 8.6 [341(wurtzite)

1400 (Loffe database)

1700

2973

AIN [30] 6.066 2.59 2030 49

1150

3273

GaN [241 3.47 0.59 600 47

600

InN [311 0.69 0.41 454 29

660

1373

Optical parameters taken from the literature have been listed in Table 1. It can be seen from Table 1 that die optical parameters have values in the decreasing order of dieir molecular size.

Application of Penn Model

For AIN, Neff = 6 electrons per atom and Perm gap calculated using equation (8) is 9.6 eV. For GaN and InN, the calculated Penn gaps are 8.2eV and 5.6eV respectively which are comparable to the corresponding peaks observed in the E2 -E spectra. All these calculated values are listed in Table 2:

Table 2: Parameters calculated using Penn model Property Density p(g/cc) Mol. Weight (W) Ncff

Au)p(eV) EF (eV) C Q O

Ep(eV)

Er. (eV)

AIN 3.26 40.99 6 19.90 15.89 4.68 9.6 6.1

GaN 6.15 83.73 6 19.12 15.07 5.7 8.22 3.5

InN 6.81 128.82 6 16.22 12.09 8.4 5.62 0.8

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Comparison of structure-related dielectric functions

We compare the structure-related A1N dielectric functions by various authors in Figure 5. The Wurtzite A1N thin films cited in this figure by Wethkamp and Wilmers[35] were grown by Metal-Organic Chemical-Vapor Phase Deposition(MOCVD) on C-AI2O3 substrates. The dielectric functions were extracted from the measured data using a spectroscopic ellipsometer operating with synchrotron radiation by an EMA correction for the surface roughness. Dominant peaks in the e2 — E spectra are at 6.13 eV, 7.76 eV and 8.79 eV. The first peak is assigned as the fundamental band gap of A1N films. The maximum at 7.76 eV is related to interband transitions at M point of the Brillouin Zone and the 8.79 eV is due to transitions at a quarter of the brillouin Zone in the I — R direction[36]. Theoretical calculations[15,37] on two wurtzite structures using modified Adachi's model dielectric function(MDF) and first-principles orthogonalized linear-combination-of-atomic orbitals method in local density approximation are introduced for comparison. The energy positions of these maxima (Table 3) are very close to the experimental data. Modified MDF seems to be closer to the experimental results than LDA first-principle calculations. For Zincblende A1N films, we take Roppischer's data[38] as an example. The sample was grown by RF plasma-assisted MBE on free 3C-SiC(001) substrates and the dielectric functions were recorded by a commercial ellipsometer. The first peak at 5.88 eV is related to the band gap. The sharp onset above 7.20 eV is correlated with a critical point around the X-point of the BZ and above 7.95 eV is due to transitions along the [11 l]-direction. Other two transitions at 11.14 eV and 12.52 eV are assigned to L- andT- points of BZ [38], For comparison, ab-initio density functional theory [39] calculated results give a slightly larger band gap (6.18 eV). Instead of a main peak at 7.20 eV, a wide peak in the range of 7 to 10 eV is observed with a maximum at 8.04 eV. One peak at 11.78 eV was present in the experimental spectrum. Finally, we present the experimental results of a single crystal samples measured by Loughin and French[40]. This single crystal was grown by a modified Bridgman technique. To obtain the dielectric functions, a laser plasma sourced VUV spectrophotometer was used for the reflectance data and Kramers-Kronig relation was applied. The e2 -E spectra shows a significant difference from those of thin films. The band gap is reported at 6.20 eV. Small peaks lower than the band gap are related to vacancies. A sharp peak is found at 8.69 eV and a smaller peak at 13.14 eV. Comparison in the experimental data between zincblende and wurtzite reveals that zincblende films starts to show non-trivial values of E2 from 5.4 eV while Wurtzite starts at around 6 eV.

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Fig, 5: Comparison of the dielectric functions of various AIN structures. Experimental results include one wurtzite thin film[35], one zincblende thin film[38] and one wurtzite single crystal[40] structure. Theoretical calculations include two wurtzite[15,37] and one zincblende[39] structures.

Fig. 6: Comparison of the dielectric functions of different GaN structures. Experimental results include two wurtzite thin films[41,42] and one zincblende thin film[42]. Theoretical calculations include one wurtzite[41] and one zincblende[43] structures.

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Fig. 7: Comparison of the dielectric functions of different InN structures. Experimental results include one wurtzite thin film[44], one ztncblende thin film[45,46] and one wurtzite single crystal[17] structure. Theoretical calculations include two wurtzite[47,48] and one zincblende[48] structures.

The experimental and theoretical spectra of GaN are put together for comparison (Fig.6). The wurtzite GaN thin film result is taken from Lambrecht and Segall[41]. They grew epitaxial layers on sapphire substrate using MOCVD. An ultrahigh-vacuum reflectometer was operated to record the reflectance data and Kramers-Kronig analysis was performed to yield the dielectric functions. The cutoff energy shows the experimental band gap of 3.4 eV. Other peaks are listed in (Table 4). These authors also calculated the theoretical curves using density functional theory(DFT) with local density approximation(LDA). Theory matches well with experiments except that the theoretical peak intensities are small and a minor peak at 19.5 eV is absent. Another measurement in the range 5 - 9.5 eV by Logothetidis[42] is also reported in Figure 6. They grew both zincblende and wurtzite GaN thin films by electron-cyclotron resonance microwave plasma-assisted MBE on silicon and sapphire substrates respectively. The dielectric functions are studied with spectroscopic ellipsometry. Based on the literature, they assigned Critical Points(CPs) at 3.35(band gap), 7.0, 7.9, 9.0 eV to T6 - r1 ; M4 - M1( M2 - Uv K3 - K2 respectively. In the mean time, they report the main features of zincblende GaN thin films at 3.25(band gap), 7.0, 7.6 eV and assign them to transitions T15 - T,, L6 - Lj. and X6 - Xj. From this result, wurtzite GaN thin film has a larger band gap than zincblende GaN thin film and the former shows one more transition above 9.0 eV. Theoretical Calculations[43] of zincblende structure based on empirical pseudopotential band structures is also brought into comparison. The work reports a band gap of 3.26 eV and two peaks at 7.0 eV and 7.6 eV. These features are in good agreement with Logothetidis' results. However, the theoretical work also reports strong peaks at 8.6 eV and 12.6 eV which are not found in the experiment. The author attributed the 8.6 eV peak to transitions in the region around 7t/2a(3,l.l) in the Brillouin Zone.

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Table 3: Critical Points(band gap and peaks) in AIN e2-E spectra with different structures by various authors. Struct u-re

Techn-ique CPs

Ref.

Thin Film(Experiment) Zincblende

Spectroscopic Ellipsometry

5.88 7.20 7.95 11.14 12.52

38

Wurtzite Spectroscopic Ellipsometry

6.13 7.76 8.79

35

Single Crystal Experiment

VUV Spectrometer

5.12 6.20 6.80 8.68 13.14

40

Theory Wurtzite

First principle LCAOLDA

4.94 6.74 7.50 9.23 10.26 11.84 14.82

37

Wurtzite M-

MDF 6.00 6.68 7.60 8.96

15

Zincblende Ab initio DFT+QP

5.86 6.18 7.12 8.04 9.40 11.78

39

Table 4: Critical Points(band gap and peaks) in GaN E2-E spectra with different structures by various authors. Structure

Techniqu-es

CPs

Ref.

Thin Film(Experimenf) Zincblende

Spectroscopic Ellipsometry

3.25 7.00 7.60

42

Wurtzite KKR+UV

Reflectometer 3.40 6.96 7.82 9.23 10.53 12.65 19.5

41

Wurtzite Spectroscopic Ellipsometry

3.35 7.00 7.90 9.00

42

Theory Zincblende Empirical

pseudopotential 3.26 7.00 7.60 8.60 12.60

43

Wurtzite DFT+LDA

3.40 7.00 7.98 9.23 10.53 12.38 12.86 13.30 41

Table 5: Critical Points(band gap and peaks) in InN E2-E spectra with different structures by various authors.

Structure

Techniq-ues

CPs

Thin Film(Experiment) Zincblende

0.59 1.27 5.12 6.17 9.28

Wurtzite Spectroscopic Ellipsometry

0.76 4.85 5.38 5.80 6.18

Single Crystal Experiment SOR-RING

Monochrometer KKR 1.90 2.36 4.77 7.66 8.52

Theory Zincblende FPLAPW

1.03 5.47 6.92 9.20 11.43 11.72 16.33

Wurtzite Ab initio DFT+LD

A/QP 0.93 4.62 6.62 8.22 8.62 9.48 10.43 12.38

Wurtzite FPLAPW

1.39 1.82 6.00 8.47 9.28 10.22 11.22 11.67

710

Ref. 45,46 44 17 48 47

12.93 13.93 16.28

48

Table 6: List of experimental structure, thickness, band gap of A1N, GaN, InN by various authors.

Material

A1N

GaN

InN

Structure Wurtzite thin film

Zincblende thin film Wurtzite Single crystal Wurtzite thin film(l) Wurtzite thin film(2) Zincblende thin film

Wurtzite thin film Zincblende thin film

Wurtzite Single crystal

Thickness(nm) 150

40, 100 -

45 4000 700

700,950 -

200-300

Band gap(eV) 6.13 5.88 6.20 3.40 3.35 3.25 0.76 0.59 1.90

Reference 35 38 40 41 42 42 44

45,46 17

The comparison of experimental and theoretical dielectric functions of InN with respect to different structures are shown in Figure 7. We utilize the spectroscopic ellipsometry data from Maria Losurdo and Giovanni Bruno[45] who deposited InN films on Si face of 6H-SiC substrates by MBE. The fundamental band gap is found at 0.76 eV. There are also other four main features listed in Table 5 and the corresponding electronic transitions are summarized by the author[44]. For zincblende InN thin films, we take the spectrum reported by Goldhahn[45,46], The critical points of the experiment are recorded at 0.59(band gap), 1.27, 5.12, 6.17, 9.28 eV. The band gap is much smaller than that of the wurtzite structure. A number of theoretical work has been done on InN. One calculation[47] using ab initio method combined with DFT-LDA and quasiparticle gives band gap of 0.93 eV. Another calculation by Persson et al[48] using a full-potential linearized augmented plane wave method gives a band gap of 1.39 eV. Persson et. al also reported their calculation on InN with zincblende structure. A band gap of 1.03 eV is derived and the critical points in the e2 -E spectra match very well with experiments. [46]. More peaks higher than 10 eV are found in the calculations. Finally, we report a single crystal InN spectrum by Guo et. al[17]. The single crystal films were grown on (0001) a-AI2O3 substrates by microwave-excited MOVPE. Reflectance data was measured at the first line of the SOR-RING with a 1-m Seya-Namioka type monochrometer. Dielectric functions are obtained by applying Kramers-Kronig relations to the reflectance spectra. The main features of the s2 -E spectra are above 2.36, 4.77, 7.66, 8.52 eV. The lower energy spectrum shows an essential agreement with other calculations while the higher energy part gives great discrepancy.

Conclusions

Optical properties of A1N, GaN and InN reported by various authors are compared and analyzed. We apply the Penn model to all three materials to study the optical properties. Temperature and pressure dependent band gap are plotted and discussed. Finally, we review the structure-dependent dielectric functions from both experimental and theoretical perspectives.

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TIMIS2012 141 s t Annual Meeting & Exhibition

Science and Engineering of Light Metal Matrix Nanocomposites and

Composites

Edited by: Xiaochun Li

Alan Luo

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

INTERFACIAL ANALYSIS OF CNT REINFORCED AZ61 MG ALLOY COMPOSITES

Katsuyoshi Kondoh', Hiroyuki Fukuda2, Junko Umeda', and Bunshi Fugetsu3

1 JWRI, Osaka University; 11-1, Mihogaoka; Ibaragi, Osaka 567-0047, Japan 2 Graduate School of Engineering, Osaka University; 1-1, Yamadaoka; Suita, Osaka 565-0871, Japan

3 Hokkaido University; Niow5, Kita-ku; Sapporo, Hokkaido 060-0810, Japan

Keywords: Magnesium, Carbon nanotube, Carbide, Interface, Powder metallurgy

Abstract

Magnesium (Mg) composite reinforced with carbon nanotubes (CNTs) having superior mechanical properties was fabricated by using both pure Mg and AZ61 Mg alloy matrix in this study. The composites were produced via powder metallurgy route containing wet process using isopropyl alcohol (IPA) based zwitterionic surfactant solution with unbundled CNTs. The produced composites were evaluated with tensile test and Vickers hardness test and analyzed by X-ray diffraction (XRD) and field-emission scanning electron microscopy (FE-SEM) equipped with energy dispersive spectroscopy (EDS) and electron back scattered diffraction (EBSD). As a result, only with AZ61 Mg alloy matrix, tensile strength of the composite was improved. In-situ formed A12MgC2 compounds at the interface between Mg matrix and CNTs effectively reinforced the interfacial bonding and enabled tensile loading transfer from the Mg matrix to nanotubes. Furthermore, it was clarified that the microstructures and grain orientations of the composite matrix were not significantly influenced by CNT addition.

Introduction

Recently, due to the demand of superior fuel consumption and subsequent lower CO2 gas emission, the application fields of lightweight materials are expanding wider and greater. In particular, magnesium (Mg), which is the lightest structural metal with the density of 1.74 g/cm3, and its alloys are one of the most promising materials. However, their mechanical properties are weak compared to other structural metals such as aluminum (Al) and titanium (Ti) alloys. For instance, ultimate tensile strength of the mould casting AZ91C Mg alloy is only 275 MPa [1] while that of the casting AZ6N4 Al alloy is 540 MPa [2]. In spite of the utilization demand of lightweight materials, their inferior mechanical strength prevents the wide application of Mg and its alloys at present. Thus, many attempts to enhance their mechanical properties are conducted in lots of ways such as, famously, severe plastic deformation processing for grain refinement [3] and incorporation of reinforcements [4]. On the other hand, carbon nanotubes (CNTs) having cylindrical shape of the scrolled graphen sheets [5] exhibit remarkable mechanical properties. According to the reported article [6], their Young's modulus is approximately 4 TPa. However, since strong Van der waals forces are induced between CNTs due to thier nano-order size, CNTs easily agglomerate into the large bundles up to several hundred microns. As easily understood, when such bundled CNTs are incorporated into matrix as reinforcements, they behave as defects

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in the composite matrix and result in decrease of the composite properties. Nonetheless, in the previous article, authors succeeded in fabricating high-strength pure Mg and AZ31B Mg alloy matrix composites reinforced with CNTs by employing their original wet process using powder metallurgy route [7]. By immersing into the water based solution containing zwittenonic surfactant having both positive and negative charges at the tip of its hydrophilic group, CNT bundles were effectively disassembled [8]; thus, CNTs could be successfully distributed in the Mg matrix by employing this technique. However, the excess amount of MgO synthesized during the composite fabrication process caused much reduction of the ductility of the produced composites whereas some of MgO strengthened the interface between CNT and Mg matrix. In contrast, Kwon et al. produced pure Al composite with CNTs having excellent strength without losing its elongation [9]. Despite the high strength improvement, their composite maintained more than 10 % true strain. According to their analysis, the interface between Al matrix and CNTs was mainly reinforced by in-situ formed aluminum carbide (AI4C3) discontinuously distributed at the structural defects of CNTs without introducing destructive damage to CNTs. In this point of view, the production key of MMCs with CNTs having high strength with acceptable ductility is to effectively reinforce the interface by synthesizing carbides. Nevertheless, elemental Mg and carbon (C) can not easily produce magnesium carbides as widely recognized. One of authors clarified that Mg and C did not react even when they were thermally deposited together in atomic state in his earlier work [10]. Thus, it is difficult to reinforce the interface between Mg matrix and carbon materials by producing the binary magnesium carbides. However, a few reports discovered a formation of the ternary carbides of Al2MgC2 [11-16]. Viala et al. concluded that, in Mg matrix containing more than 0.6 wt % up to 19 wt % Al, Mg, Al and C formed ternary carbides of Al2MgC2 at the interface between carbon fiber and Mg matrix [11]. Besides, Pei et al. observed that tensile strength of AZ91 Mg alloy was improved by synthesizing Al2MgC2 at the interface between me Mg matrix and carbon fibers [12], Hence, in the present study, firstly, pure Mg composite reinforced with CNTs (pure Mg/CNT composite) was produced by using the isopropyl alcohol (1PA) based zwitterionic surfactant solution containing CNTs instead of the water based solution employed in the previous study to understand the effectiveness of MgO prevention on the mechanical properties [7, 17]. Employing the IPA based solution was capable of avoiding the excess amount of MgO formation during sample fabrication steps. Then, AZ61 Mg alloy composite with CNTs (AZ61/CNT composite) was also produced. By using AZ61 Mg alloy containing 6 wt % Al, Al2MgC2 ternary carbide was synthesized at the interface and the carbide strengthened the interfacial bonding between Mg matrix and CNTs. The produced composites were evaluated with tensile test and Vickers hardness. X-ray diffraction (XRD), field emission - scanning electron microscope (FE-SEM) equipped with energy dispersive spectroscopy (EDS), and electron back scattered diffraction (EBSD) system were operated to analyze the fabricated composite.

Experimental

Fabrication of Mg matrix composites reinforced with CNTs by powder metallurgy route

Pure Mg (purity; 99.5 %) and AZ61 Mg alloy powders were prepared as raw powders shown in Fig. 1 (a) and (b). The chemical compositions of AZ61 powder were Al, 6.38; Zn, 0.68; Mn, 0.28; Si, 0.04 in mass %. Multi-walled carbon nanotubes commercially named NC7000, which was the product of Nanocyl CO. LTD, were also employed in this study. From die data sheet distributed by the corporation, their mean diameter, mean length, and carbon purity were 9.5 urn, 1.5 urn, and approximately 90 %, respectively. Purchased CNTs were dispersed in the IPA based zwitterionic surfactant solution to disassemble their bundles. The disassemble mechanism of the

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bundled CNTs were described in another article [8]. Then, the solution was mixed with as-received powders to transfer CNTs onto the powder surface. The paste-like composite powders with CNTs were dried at 80 °C for 3 hours and heated at 500 °C for 1 hour in argon (Ar) gas atmosphere to thermally decompose the residual surfactant elements from the powder surface (debinding process). In the prior studies, the water based zwitterionic surfactant solution was used to disassemble CNT bundles [7, 17]. However, since water reacted with Mg powder surface and provided MgO, which highly decreased the composite elongation, for the purpose of preventing MgO production, the IPA based solution was employed in the present study. After removing the residual surfactants, the composite powders were consolidated into columnar billets having 42 mm of diameter and about 25 mm of height with spark plasma sintering (SPS) [18] system at 550 °C for 30 minutes with 30 MPa pressure in vacuum. Further, to remove the pores introduced during SPS processing, hot extrusion was applied to the SPSed billets at 400 GC for pure Mg based ones and 350 °C for AZ61 based ones at the extrusion speed of 0.3 mm/s. Before the hot extrusion, the billets were kept at the above temperatures for 5 minutes in Ar gas atmosphere. The extruded rods had 7 mm in diameter and about 1 m in length. With these extruded rods, C content of each sample was measured by the infrared-absorption technique; the measured C contents were 0.26 vol % for pure Mg/CNT composite and 0.74 vol % for AZ61/CNT composite. The measured C contents were regarded as CNT contents of the composites because no other carbon source was included in these composites. Besides, some portion of the extruded rods were subsequently heated at 550 °C for 10 hours under Ar gas atmosphere to facilitate Al2MgC2 formation since its formation needs long heating at least close to melting point of Mg according to references [11, 13]. ,.. , ,,c c c x . , ,,

° ° Fig. 1 FE-SEM observation on as-recieved pure „ , ^ r . *_ . i . t. • i Mg powder (a) and AZ61 Mg alloy powder (b). Evaluation of microstructural and mechanical properties of Mg alloy composites

In this study, extruded rods were evaluated with tensile test and Vickers hardness test. For the tensile test, tensile loading system was operated with the strain rate of 5.0 * 10"4 /s. Rod type specimens having 15 mm in gauge length and 3mm in diameter were prepared for the test. Vickers hardness test was conducted with loading weight of 5.0* 10"2 kfg for 15 seconds on both pure Mg and AZ61 based composites. Microstructures were characterized with XRD, FE-SEM equipped with EDS, and EBSD. Before observations, the extruded composites were sectioned parallel to the extrusion direction, abraded with #600, 1500, and 4000 SiC abrasive papers, and then, buffed with 0.25 urn diamond paste. After each step, specimen surfaces were carefully rinsed in ethanol with ultrasonic vibration. For EBSD analysis, the buffed samples were electrochemically finished with the mixture of 150 ml of phosphoric acid and 250 ml of ethanol at the voltage of 5 V for 30 seconds. After the electrochemical finishing, the specimen surfaces were cleaned with methanol and subsequently washed in ethanol with ultrasonic vibration to completely remove the surface impurities.

Results and discussion

Oxide contamination in the composites

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XRD profiles of pure Mg/CNT composite powders are indicates the composite powders produced with the IPA based surfactant solution after debinding, that before debinding, as-received pure Mg powders, and the composite powders produced with the water based surfactant solution after debinding process in reference [7], respectively. In only (d), peaks of MgO were identified though other profiles of (a), (b), and (c) showed only peaks of Mg. In the previous study, Mg(OH)2 and MgO were identified in the composite powders before and after debinding process, respectively [7]. Authors explained that, during debinding process at 500 °C, Mg(OH)2 was dehydrated into MgO. This explanation was in accordance with the dehydration temperature of Mg(OH)2 at 300 °C [19]. However, in the present study, neither Mg(OH)2 in the composite powders before debinding of (b) nor MgO in that after debinding of (a) were identified in the XRD profiles; the profiles showed only Mg peaks as it was in as-received pure Mg powders of (c).

presented in Fig.2. (a), (b), (c), and (d)

Fig.2 XRD profiles of pure Mg/CNT composite fabricated with IPA based zwitterionic surfactant solution containing CNTs after (a), before the debinding process (b), as-received pure Mg powders (c), and the composite powders produced with the water based solution after the debinding process in reference [7] (d).

Mechanical properties and heat treatment effect of AZ61/CNT composite

Fig.3 shows nominal stress and nominal strain curves of AZ61/CNT composite before and after the heat treatment at 550 °C for 10 hours. Yield stress at 0.2 % offset and tensile strength of AZ61 Mg alloy were improved 11.05 and 7.86 MPa before the heat treatment and 10.82 and 1.82 MPa after the heat treatment by CNT addition, respectively. Both the yield stress and the tensile strength were enhanced with keeping more than 9 % elongation. Using the IPA based zwitterionic surfactant solution was effective to maintain the suitable elongation of AZ61/CNT composite. The XRD profiles of all samples are exhibited in Fig.4.

Fig.3 Tensile behavior of AZ61 with 0.74 vol% and with no CNT before (a-1) and (b-1), and after (a-2) and (b-2) heat treatment at550°Cfor!0hrs.

Fig.4 XRD profiles of AZ61/CNT composite and without CNTs before (a-1) and (b-1), and after (a-2) and (b-2) H.T at 550 °C for 10 hrs.

Peaks for Mg matrix, intermetallic compounds, AlnMn4 and Al^Mgn, and MgO were identified in all samples. MgO identified here was not produced during sample fabrication

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process but the one produced spontaneously on the as-received AZ61 Mg alloy powder surface. This is because the peak intensities of MgO were mostly same regardless of samples with or without CNTs. Therefore, MgO production was prevented as low as possible in the composite as well. However, around 2 0 = 54.7 °, an unknown peak was identified only in samples with CNTs; a small peak for the sample before the heat treatment and clear broad peak for the sample after the heat treatment. Fig.5 shows EDS elemental mapping of AZ61/CNT composite after the heat treatment. Al concentrated areas clearly surrounded CNTs. There, C which was probably diffused from CNTs as well as Mg which was concentrated less than that inside the matrix was also detected. EDS elemental mapping of AZ61/CNT composite before the heat treatment was also presented in Fig.6. In this mapping, the Al concentration and the C diffusion around CNTs were ambiguous.

Fig.5 SEM-EDS mapping on AZ61/CNT composite after H.T at 550 °C for 10 hrs.

Fig.6 SEM-EDS mapping on AZ61/CNT composite before H.T at 550 °C for 10 hrs.

However, EDS line scanning clearly showed them around the interface between Mg matrix and CNTs as shown in Fig.7. Obviously, CNTs were surrounded by the area where all Al, C, and Mg were present. Several reports identified the ternary carbides of Al2MgC2 in Mg-Al alloys containing carbon sources [11-16], One of these articles distinguished two types of Al2MgC2, Ti and T2 [11]. According to the X-ray diffraction lines of Al2MgC2 - T2 indexed in the article, the lattice spacing for the third peak of Al2MgC2 - T2 was 1.6887 Â. This can be calculated into 2 8 = 54.35 °, which was quite close to that of the unknown peak of 2 8 = 54.7 ° observed in AZ61/CNT composite both before and after the heat treatment shown in Fig.4. Here, the first and second peaks of AbMgC2 - T2 were hidden behind the larger peaks of Mg matrix and Al i2Mg|7, respectively. Besides, the direct growth of Al2MgC2 from the surface of the high strength carbon fiber was observed by Feldhoffet al. [13]. Therefore, it is reasonable that the unknown peak was derived from Al2MgC2 formed at the interface between Mg matrix and CNTs in AZ61/CNT composites both before and after the heat treatment. Since the interfacial Al2MgC2 facilitated loading transfer from the Mg matrix to CNTs, the mechanical strength of

Fig.7 SEM-EDS line scanning on AZ61/CNT composite before heat treatment at 550 °C for 10 hrs.

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AZ61/CNT composites was improved. While, before the heat treatment, AhMgCî production was less than that after the heat treatment, that amount of Al2MgC2 at the interface was enough for the interfacial bonding between Mg matrix and CNTs of the composite. In fact, in the composite after the heat treatment, high Al and C concentration areas continued to deeper in the matrix; from the EDS mapping, the area width from the interface between Mg and CNTs was around 500nm. Therefore, though the amount of Al2MgC2 was increased after the heat treatment, AbMgC2 just grew into deeper in the matrix, and so newly grown AhMgC2 due to the heat treatment was not effective to strengthen the composite. In addition, though the peak for Al2MgC2 identified in the present study was slightly broadened, this was caused by the slow formation of AhMgC2 under the slid-state reaction as Viala et al. and Bosselet et al. reported [11, 14]. Slow formation of the compound produced relatively much transformation region from Mg matrix to the compound. This region having incomplete lattice spacing caused the slight shifts of the reflected X-ray.

Conclusion

Oxide contamination was greatly removed by using the IPA based zwitterionic surfactant solution. However, tensile properties of pure Mg/CNT composite produced with the solution was not improved. Naturally formed MgO was not sufficient to strengthen the interfacial bonding between Mg matrix and CNTs. However, in this composite, elongation was improved due to the elimination of excess amount of MgO formation and the decrease of induced strain in the Mg matrix. Free CNTs provided by the destruction of interfacial MgO behaved as in-situ lubricants in the composite during the hot extrusion process. This also resulted in decrease of the extrusion load and Vickers hardness values of pure Mg/CNT composite compared to those of pure Mg without CNTs. On the other hand, AZ61/CNT composite having superior tensile strength and adequate elongation was fabricated. In the composite, Al2MgC2 ternary carbide was identified. The carbide reinforced the interfacial bonding strength between the Mg matrix and CNTs; therefore, the mechanical properties of AZ61/CNT composite was improved compared to that of AZ61 without CNTs. Although CNT addition might affect microstructure and grain orientation of the Mg matrix, FE-SEM observation and EBSD analysis clarified that neither microstructure nor texture was significantly influenced by CNT addition.

References

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12. Z. L. Pei, K. Li, J. Gong, N. L. Shi, E. Elangovan, C. Sun, Journal of Materials Science 44 (2009)4124.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

BIODEGRADABILITY AND MECHANICAL PERFORMANCE OF HYDROXYAPATITE REINFORCED MAGNESIUM MATRIX

NANOCOMPOSITES

Chao Ma1, Lianyi Chen', Jiaquan Xu1, Axel Fehrenbacher', Yan Li2, Frank E. Pfefferkorn', Neil A. Duffie1, Jing Zheng2 and Xiaochun Li'

'Department of Mechanical Engineering, University of Wisconsin-Madison 1513 University Ave.; Madison, WI 53706, USA

department of Obstetrics and Gynecology, University of Wisconsin-Madison 202 S Park St.; Madison, Wl 53715, USA

Keywords: In-vitro Corrosion, Vickers Hardness, Magnesium, Nanocomposite.

Abstract

Magnesium and its alloys have gained significant attention recently as potential alternatives for biodegradable materials due to their good biodegradability, biocompatibility and mechanical properties. However, magnesium alloys tends to have high corrosion rates in biological liquid, thus presenting a potential problem if a magnesium implant/device needs to maintain the mechanical integrity for a sufficient period under physiological conditions. It is expected that nanoparticles could help address this problem. In this study, hydroxyapatite (HA) nanoparticles were used to form magnesium-based metal matrix nanocomposites (MMNC) through friction stir processing (FSP). Microstructural study shows the HA nanoparticles were well dispersed in the magnesium matrix. While FSP alone refined grain size for magnesium, the addition of HA nanoparticles was much more significant for grain refinement. In-vitro corrosion tests were conducted in simulated body fluid (SBF), and experimental results indicated that corrosion resistance of MMNC was much improved when compared to pure Mg. The microhardness of nanocomposites was improved significantly. The study suggests that magnesium nanocomposites yield great potential for enhancing both mechanical properties and corrosion resistance for biological applications.

Introduction

Some medical implants serve no further purpose as soon as they have fulfilled their mission of assisting the healing process, such as the internal fixation device used in an orthopedic surgery [1]. It is of great interest to develop implants that are biodegradable, reducing the cost of health care and the pain of a patient from the implant removal surgery. Polymers were the first biomaterials that were developed and used for biodegradable implants. However, the inferior mechanical properties of polymers limit their applications in load-bearing and tissue supporting applications. Polymer implants are also difficult to be detected and positioned by radiography due to their radiolucency. On the other hand, metals offer relatively high mechanical properties and can be readily detected by radiography. Unfortunately, most metals are either not degradable within body or toxic to human biological system. As one of the rare exceptions, magnesium and its alloys have emerged as a promising class of biomaterials. It is found that the degradation products of magnesium are able to be metabolized by the body, and thus biocompatible and biodegradable. Magnesium has similar mechanical properties to human bone and thus can serve as excellent fixation implants during orthopedic surgeries. However, the low corrosion resistance

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of magnesium, which results in its dissolution in human body much sooner than desired, is one of the major concerns for its widespread applications in biomédical field. The rapid corrosion will damage the mechanical integrity of the implants during healing process. In addition, the rapid release of corrosion products such as hydrogen gas and hydroxide could also affect the healing process [2-7].

Many techniques have been developed to improve corrosion resistance of magnesium, including surface treatment, alloying and composites. It is particularly of interest to investigate the corrosion properties of magnesium metal matrix composite (MMC) and nanocomposite (MMNC) as their properties could be controlled by varying the percentage of particles. Past investigations have been more focused on bioceramic particles reinforced magnesium MMC and MMNC through powder metallurgy (PM) technique. Hydroxyapatite (HA) is one of the most popular bioceramic reinforcements because it has similar chemical and crystallographic structures to bone and its bioactivity in body is expected. Witte et al [8] in one of the first reports on magnesium MMC as biodegradable material experimented with cytocompatible AZ91D/HA (20 wt.%) MMC by PM method. The conglomerations of HA particles were found inhomogeneous spread over the sample, and the size differed from less than 1 um to 20 urn and even more. The corrosion resistance of AZ91D/HA was improved compared with that of AZ91D matrix. However, aluminum content in AZ91D is considered as toxic to human nervous system. Therefore, it would be of interest to investigate the properties of HA reinforced pure magnesium. Gu et al [9] adopted PM method and made Mg/HA MMC (10 wt.%, 20 wt.% and 30 wt.%). Spherical HA particles of 2-3 um were used as starting enhancement material but agglomerations of the particles were observed with cluster size varying from 10 um to 20 p.m. The yield tensile strength of 10 wt. % Mg/HA composite was 9% higher than that of the pure Mg, while the yield tensile strengths of 10 wt. % and 20 wt. % Mg/HA were 3% and 33% lower than that of pure Mg. The ultimate tensile strength and ductility of Mg/HA composites decreased with the increase of HA content. The corrosion resistance of Mg/HA composites in SBF also decreased with the increase of HA content. Liu et al [10] added 20 wt. % rod-like HA particles (100 nm in length and 500 nm in diameter) into pure Mg with PM method and found the corrosion rate of Mg/HA MMNC is decreased by 30% compared with pure Mg. The mechanical performance was not reported in this paper. The contradictory results might be due to the clustering of HA particles and porosity, as it is difficult to disperse these particles well through PM method. Better processing technology and more fundamental study are needed to understand the phenomena better. Frication stir processing (FSP) has been used to fabricate composites in several studies [11-17], in which the dispersion of the particles and the interfacial bonding between the particles and the matrix were found very excellent. This paper utilizes FSP to disperse HA nanoparticles to pure magnesium for further study, especially on its microstructure, mechanical and corrosive properties.

Experimental Methods

Fabrication of Mg/HA Composites

Pure Mg (99.93%) ingot was used. The reinforcements were needlelike HA nanoparticles with an average size of 20 nm. FSP process was used to disperse HA nanoparticles into pure Mg matrix. The schematic of the fabrication process is shown in Figure 1. As shown in Figure la, two Mg plates were used. The dimension of the Mg plates was 120 mm x 50 mm x 3.5 mm. A groove of 100 mm x 1 mm (width) x 2 mm (depth) was milled onto the top plate. HA nanoparticles were filled into the groove and condensed to achieve a higher HA loading. The top

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plate filled with nanoparticles was flipped over and placed on the bottom Mg plate to prevent the nanoparticles being squeezed out of the Mg matrix during FSP. FSP was conducted to mixed the HA nanoparticles into Mg matrix after clamping the two plates on the FSP machine and aligning the travelling path of the tool with the groove. The FSP machine used in this study was a CNC milling machine with a 7.5 HP spindle. The pin of the tool was threaded and it was 5 mm in diameter and 5 mm in depth. The shoulder diameter was 15 mm. The rotation speed of the spindle was set to 1600 rpm and the travelling speed of the tool was 25 mm/min. A 3 " tilt angle was also applied. Four passes of FSP were applied back and forth to disperse the nanoparticles well. Assuming the nanoparticles move only laterally other than vertically, the weight fraction of HA in Mg/HA nanocomposite is estimated to be 15%. Two pure Mg plates without any groove (thus no nanoparticles loaded) were also welded together with the same parameters for comparison. Pure Mg plates without FSP were also studied to serve as the reference material. The denotations of these three groups are listed in Table I.

Figure 1. The cross section view of the fabrication process: a) before FSP; b) during FSP.

Table I. The denotations of the three groups of materials for comparison. Denotation N Description | as-cast pure Mg plates | FSP processed Mg Mg/HA nanocomposite

Microstructure

For microstructural study, samples of the three groups (P, F and N) were cut, mounted, mechanically ground, polished down to 1 um and finally ultrasonic cleaned in ethanol for 1 min. Specimens were then etched with a solution containing 10 mL acetic acid, 4.2 g picric acid, 10 mL deionized water and 70 mL ethanol.

The microstructures were observed by optical microscopy (OM), polarized light microscopy (PLM) and scanning electron microscopy (SEM).

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Vickers Hardness Tests

Vickers hardness was used to indicate the mechanical properties of the samples. The tests were conducted using a 200 gf load for 10 s. Five measurements on each sample were taken and the average value was used. The Vickers hardness of the three groups (P, F and N) were tested and compared.

In-vitro corrosion tests

The three groups of samples were cut to get desired dimension, 6 mm x 6 mm x 2 mm. The samples were mechanically ground with sand paper till 1200 grit, and then ultrasonic cleaned in ethanol for 5 min.

The corrosion tests were carried out in modified simulated body fluid (m-SBF) at a temperature of 37 °C with an oscillator. The chemical composition of m-SBF is closer to the ion concentrations of human blood than that of conventional simulated body fluid (c-SBF) [18].

After different immersion time, the samples were taken out from the m-SBF, rinsed with ethanol and dried. The samples were then immersed in chromic acid solution (200 g/L Cr03, lOg/L AgN03) for 10 min to remove the corrosion products before weighting. The in-vitro corrosion rates were calculated by weight loss method:

where CR is the corrosion rate, Am is the weight loss, p is the density of material being tested, A is the initial surface area and t is the immersion time. For each type of material, two samples were tested at each time and the average results were used.

Results

Microstructure

Figure 2 shows the SEM images of the cross section of a representative Mg/HA nanocomposite sample under different magnifications. The bright spots are HA nanoparticles in the dark Mg matrix. No large clusters of the HA nanoparticles were observed. It also shows the uniform distribution of the HA nanoparticles in Mg matrix.

Figure 3 presents PLM images of the cross section of the three groups of samples respectively. The pure Mg in Figure 3a shows three large grains (grain boundaries shown as blacked curves) with subgrains. In Figures 3b and 3c, different colors indicate different grains because the colors also change differently when the samples were rotated. The grains of Mg are much refined after FSP (Figure 3b) because of the dynamic recrystallization induced by significant frictional heating and intense plastic deformation during FSP [19]. Figure 3c shows that adding HA nanoparticles into Mg matrix refines the grain much more significantly.

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Figure 2. The SEM images of the cross section of Mg/HA nanocomposite sample.

Figure 3. The PLM images of the cross section of the samples: a) as-cast pure Mg; b) FSP processed Mg and c) Mg/HA nanocomposite.

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Vickers Hardness Tests The comparison of Vickers hardness of the three groups of samples is shown in Figure 4. The hardness of the FSP processed sample is 19% higher than that of the as-cast one while the hardness of the Mg/HA MMNC is 68% higher than that of the FSP processed pure Mg. The significant enhancement in the hardness of Mg/HA MMNC is not only because of grain refinement but also contributed by second phase strengthening of HA nanoparticles.

Figure 4. The comparison of Vickers hardness.

In-vitro corrosion tests

Figure 5 represents the corrosion rates of the three groups of materials. The results show that FSP process helped reducing the corrosion rate for pure Mg. It is clear that the corrosion rate of Mg/HA is the lowest among the three groups of samples. The enhancement of corrosion resistance could be attributed to grain refinement and the effects of HA nanoparticles.

Figure 5. The comparison of corrosion rates.

The grain refinement helped to improve the resistance to pitting corrosion, which was the main corrosion mechanism of pure Mg. It was found that metastable pits that occurred on the

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microcrystalline pure aluminum were hard to develop more deeply while the metastable pits on the coarse grained counterpart exhibited higher probability to grow into deep stable pits [20]. The same mechanism might apply for the pitting corrosion behavior of pure Mg. On the other hand, the HA particles impeded galvanic corrosion by restriction of the motion of free electrons as ceramic barriers. It was also proposed that HA could induce precipitation of compounds such as calcium phosphate and formation a protective layer on the surface of the composite [8, 10]. The potential causes for the enhancement of corrosion resistance still need further and detailed study.

Conclusion

HA/Mg MMNC was fabricated through FSP. It was shown from microstructural study that the HA particles were uniformly distributed and well dispersed in Mg matrix. Much grain refinement was achieved by FSP alone while adding HA nanoparticles refined the grains more significantly. Vickers hardness of Mg/HA MMNC was improved greatly compared with both FSP processed and as-cast pure Mg. The corrosion resistance was also increased by FSP process and addition of HA. The grain refinement might improve the pitting resistance. The addition of HA decreased the galvanic corrosion and might also help form a protective layer.

References

1. O. Bostman and H. Pihlajamaki, "Clinical Biocompatibility of Biodegradable Orthopaedic Implants for Internal Fixation: A Review", Biomaterials, 21 (2000), 2615-2621.

2. M. P. Staiger et al., "Magnesium and its Alloys as Orthopedic Biomaterials: A Review", Biomaterials, 27 (2006), 1728-1734.

3. Zeng, R., Dietzel, W., Witte, F., Hort, N. and Blawert, C, "Progress and Challenge for Magnesium Alloys as Biomaterials", Advanced Engineering Materials, 10 (2008), B3-B14.

4. H. S. Brar, M. O. Platt, M. Sarntinoranont, et al, "Magnesium as a Biodegradable and Bioabsorbable Material for Medical Implants", JOM, 61 (2009).

5. C. K. Seal, K. Vince and M.A. Hodgson, "Biodegradable Surgical Implants Based on Magnesium Alloys - A Review of Current Research", IOP Conf. Series: Materials Science and Engineering, 4 (2009), 1-4.

6. H. S. Brar, B. G. Keselowsky, Malisa Sarntinoranont, Michèle V. Manuel, "Design Considerations for Developing Biodegradable and Bioabsorbable Magnesium Implants", JOM, 63(2011).

7. Y. Zheng and X. Gu, "Research Activities of Biomédical Magnesium Alloys in China", JOM, 63(2011).

8. F. Witte, F. Feyerabend, P. Maier, J. Fischer, M. Stornier, C. Blawert, W. Dietzel, N. Hort, "Biodegradable Magnesium-hydroxyapatite Metal Matrix Composites", Biomaterials, 28 (2007), 2163.

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9. X. Gu et al., "Microstructure, Mechanical Property, Bio-corrosion and Cytotoxicity Evaluations of Mg-HA Composites", Mater. Sei. Eng. C, 30 (2010), 827-832.

10. D-B. Liu, M-F. Chen and X-Y. Ye, "Fabrication and Corrosion Behavior of HA-Mg-Zn Biocomposites", Front Mater. Sei. China, 4 (2010), 139-144.

11. C.I. Chang, Y.N. Wang, H.R. Pei, C.J. Lee, and J.C. Huang, "On the Hardening of Friction Stir Processed Mg-AZ31 Based Composites with 5-20% Nano-ZrCh and Nano-SiC>2 Particles", Mater. Trans., 47 (2006), 2942-49.

12. C.J. Lee, J.C. Huang, and P.J. Hsieh, "Mg Based Nano-composites Fabricated by Friction Stir Processing", Scripta Mater., 54 (2006), 1415-20.

13. C.J. Lee and J.C. Huang, "High Strain Rate Superplasticity of Mg Based Composites Fabricated by Friction Stir Processing", Mater. Trans., 47 (2006), 2773-78.

14. Y. Morisada, H. Fujii, T. Nagaoka, and M. Fukusumi, "MWCNTs AZ31 Surface Composites Fabricated by Friction Stir Processing", Mater. Sei. Eng., A, A419 (2006), 344-48.

15. Y. Morisada, H. Fujii, T. Nagaoka, and M. Fukusumi, "Effect of Friction Stir Processing with SiC Particles on Microstructure and Hardness of AZ31", Mater. Sei. Eng., A, A433 (2006), 50-54.

16. M. Dixit, J.W. Newkirk, and R.S. Mishra, "Properties of Friction Stir-processed Al 1100-NiTi Composite", Scripta Mater., 56 (2007), 541-44.

17. Wei Wang, Qing-yu Shi, Peng Liu, Hong-ke Li, Ting Li, "A Novel Way to Produce Bulk SiCp Reinforced Aluminum Metal Matrix Composites by Friction Stir Processing", Journal of Materials Processing Technology, 209 (2009), 2099-2103.

18. Oyane A, Kim H, Furuya T, Kokubo T, Miyazaki T, Nakamura T, "Preparation and Assessment of Revised Simulated Body Fluids", J Biomédical Mater Res A, 65 (2003), 188-195.

19. Z.Y. MA, "Friction Stir Processing Technology: A Review", Metallurgical and Materials Transactions, 39A (2008), 642-658.

20. Guozhe Meng, Liyan Wei, Tao Zhang, Yawei Shao, Fuhui Wang, Chaofang Dong, Xiaogang Li, "Effect of Microcrystallization on Pitting Corrosion of Pure Aluminum", Corrosion Science, 51 (2009), 2151-2157.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

MECHANICAL PROPERTIES OF A356-CNT CAST NANO COMPOSITE PRODUCED BY A SPECIAL COMPOCASTING ROUTE

Benyamin Abbasipour, Behzad Niroumand, Sayed Mahmoud Monirvaghefï Department of Materials Engineering, Isfahan University of Technology, Isfahan, 84156-83111,

Iran

Keywords: A356-CNT nano composite, Ni-P electroless plating, Stir casting, Compocasting, Microstructure, Mechanical properties, High temperature tensile strength.

Abstract A356 aluminum alloys reinforced with carbon nano-tubes (CNTs) were produced by stir casting and compocasting routes and their microstructural characteristics and mechanical properties in ambient and high temperature were examined. In order to alleviate the problems associated with poor wettability, agglomeration and gravity segregation of CNTs in the melt, CNTs were introduced into the melts by injection of CNT deposited aluminum particles instead of raw CNTs. The slurries were subsequently cast at temperatures corresponding to fully liquid as well as 0.15 and 0.30 solid fractions. The results showed that addition of CNTs to A356 matrix could significantly increase the mechanical properties. Yield strength and ultimate tensile strength of the composite samples were significantly improved compared with the corresponding monolithic samples. While the monolithic sample kept about 60% of its strength at 300°C, the semi-solid composite samples kept more than 90% of their strength at this temperature.

1. Introduction Metal matrix composites (MMCs) are an important group of structural materials used in automotive, defense and aerospace applications because of their low density, high specific strength and modulus, excellent wear resistance, higher service temperature and in general better physical and mechanical properties compared with the corresponding monolithic materials [1]. Among a variety of MMCs produced in last few decades, aluminum matrix composites reinforced with various particles have attracted many researchers. In recent years, nano-sized particles have also drawn much interest as reinforcements in MMCs because of their superior properties compared with those of micro-sized particles [2-4], Since the discovery of carbon nano-tubes (CNTs) and their novel properties by IIJIMA [5], many potential applications have been proposed and investigated for CNTs [6]. The high strength, elastic modulus, flexibility and unique conductivity of CNTs [7] along with many other fascinating properties have led to their use in a variety of nano-composite materials. Although various investigations on fabrication of Al-CNT nano-composite have been reported in recent years, so far most of the CNT reinforced composites have been produced by solid state processes-|Ml,] 8which are not apt for production of large and complicated components. Although liquid state processes such as vortex and squeeze casting methods [12-13] provide greater freedom in component design and manufacturing, very large specific surface area and high interfacial energy of CNTs result in their agglomeration and poor distribution in the melt. Therefore, special techniques are required for addition of CNTs to the melt. Compocasting is a liquid state process in which the reinforcement particles are added to a solidifying melt while being vigorously agitated. It has been shown that the primary solid particles already formed in the semi-solid slurry can

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mechanically entrap the reinforcing particles, prevent their gravity segregation and reduce their agglomeration [14-16], These will result in better distribution of the reinforcement particles. The lower porosity observed in the castings has been attributed to the better wettability between the matrix and the reinforcement particles as well as the lower volume shrinkage of the matrix alloy. In a previous paper [17], the authors introduced a special method for addition of CNTs to semi-solid metal slurries which would result in less agglomeration and better distribution of CNTs in the matrix. It was shown that addition of CNTs to A356 aluminum alloy with this method, could refine the microstructure and improve the hardness of the cast nano composites. In this study, other mechanical properties of the cast nano composites including the yield strength, ultimate tensile strength and elongation in ambient and high temperature are examined.

2. Experimental procedure A356-CNT composites were produced by a special compocasting method using A356 aluminum alloy as the matrix and multi-walled carbon nano-tubes (MWCNTs) as the reinforcement. MWCNTs with 95% purity and outer diameters in the range of 10-30 nm were acquired from Research Institute of Petroleum Industry (RIPI) in Tehran, Iran. A transmission electron microscopy (TEM) image of the as-received MWCNTs is shown in Fig (1).

Figure 1. TEM micrograph of as-received MWCNTs.

First, CNTs were co-deposited with Ni and P as a composite layer on aluminum particles of commercial purity using Ni-P electroless plating technique. The details of the electroless plating process in described elsewhere [17]. Then, the CNT deposited aluminum particles were injected into the melt at a temperature above its liquidus temperature (700 CC) while being isothermally stirred. lwt%Mg was added to the melt to increase the wettability between the matrix and the reinforcements. After completion of the injection, the slurry was cast into a steel die placed below the furnace, either at 700 °C or after continuously cooling to 610 °C and 601 °C, corresponding to 0.15 and 0.30 solid fractions according to Scheil equation, respectively. Analogous castings but without reinforcement addition were also produced for comparison. Table I shows the experimental conditions used in different experiments.

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Table I, Experimental conditions used in different experiments, „ . CNTs Reinforcement Casting temp. Stirring speed

P (Vol%) injection temp. (°C) (°C) (rpm) A-0-0 A-15-0 A-30-0 A-0-2 A-15-2 A-30-2

0 0 0 2 2 2

700 700 700 700 700 700

700 610 601 700 610 601

500 500 500 500 500 500

The cast samples were sectioned and their microstructural characteristics were studied using an optical microscope (OM) connected to an image analysis system and a Philips XL30 scanning electron microscope (SEM). Deep etching was also carried out on the surface of one of the composite (sample A-30-2) to validate the presence of CNTs in the matrix of the cast composite. Low and high tensile tests were performed on the samples according to ASTM B557M standard. Researches have shown that the strength of A356 aluminum alloy decreases significantly at temperatures above 0.54Ts, where Ts is the solidus temperature of the alloy [18, 19]. Therefore, 300 °C was chosen for high temperature tensile tests.

3. Results and discussion Figs 2(a) and 2(b) show the optical microstructures of the monolithic A356 aluminum alloy cast at 0 and 30% solid fractions, respectively (samples A-0-0 and A-30-0). Higher magnification micrographs are also shown as inlays.

Figure 2. Typical micrographs of cast samples: (a) A-0-0 and (b) A-30-0.

Dendritic microstructures of the samples cast from full liquid state and non-dendritic microstructures of those cast from within the semi-solid temperature range are clearly visible. The microstructure of the semi-solid samples as shown in Fig 2(b) contains the primary particles solidified before pouring, i.e. during stirring and slow cooling stage, as well as the secondary particles solidified at a higher cooling rate in the steel die and the eutectic constituents. Fig 3 shows the deep-etched surface of the composite sample A-30-2 and the tips of the CNTs protruded from the composite surface due to deep etching of the matrix. As can be seen, by using

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this special technique for producing A356-CNT composites, the CNTs have been distributed uniformly in the matrix and a limited degree of agglomeration is observed which is expected to result in better mechanical properties of the cast composites.

Figure 3. Deep etched surface of the composite sample A-30-2.

Fig 4 illustrates the stress-strain curve for the monolithic sample A-0-0 produced by melt stirring and for the composite sample A-30-2 produced by the special compocasting route. As can be seen, the mechanical properties such as yield strength, elastic modulus, ultimate tensile strength and the toughness of the composite sample has significantly increased by addition of CNTs and using the advantages of semi-solid casting route for producing the composite. In fact, mechanical properties of all the composite samples were increased compared with the corresponding monolithic samples.

Figure 4. Stress-strain curve for the samples A-0-0 and A-30-2.

The mechanical properties of all the casting samples are reported in Table II. The data reported in the table is the average of three measurements. The increase in the mechanical properties can partly be attributed to the decrease in the grain size of the matrices of the composite samples. Furthermore, like other reinforcements, CNTs strengthen the matrix by creating a high density of dislocations during cooling to room temperature due to the difference of the coefficients of

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thermal expansion between the CNTs and the matrix. Mismatch strains developed at the interfaces of CNTs and the matrix obstruct the movement of the dislocations, resulting in the improvement of the mechanical properties. Improvement in the tensile strength and elongation with only 2 vol.% CNT addition indicates a resistance to plastic deformation in the stressed state due to difficulty of the dislocations rearrangement.

Table II. Mechanical properties of the cast samples. Sample Ys (MPa) UTS (MPa) Elongation (%) A-0-0 A-0-2 A-15-0 A-15-2 A-30-0 A-30-2

92 128 98 143 113 158

114 146 121 152 127 195

3.5 2.9 4.5 4.2 4.6 4.4

Table II also shows that the mechanical properties of the composite samples have been improved by casting from the semi-solid state. This is attributed to the higher density and less gas and shrinkage defects formed in the compocast samples, as shown by previous studies [14]. The 0.2% Ys of all the cast samples at two different temperatures of 25 and 300 °C are presented in Table III. Also shown in the table is the percentage of the retained yield strength at high temperature of the samples.

Table III. 0.2% Ys of the cast samples at 25 and 300 °C as well as the percentage of the retained yield strength at 300 °C.

„ . Ysat25°C Ysat300°C Retained yield strength sample ( M p a ) (MPa) at 300 °C (%)

60 82 71 87 71 90

As it is evident from the table, the strength of all the samples has decreased at higher temperature. This can be attributed to the softening the matrix alloy at high temperature. However, as it can be seen from the table, while the composite sample A-30-2 has been able to maintain about 90 percent of its strength at 300 °C, the monolithic sample A-0-0 has maintained only about 60 percent of its strength at 300 °C. It seems that presence of the CNTs as reinforcement in the matrix alloy and high thermal stability of these particles prevents the plastic flow of the matrix at high temperature. The fractographes of the composite sample A-30-2 after the tensile test at 25 °C and 300 °C are shown in Fig 5. As shown in Fig. 5a, the fracture surfaces of the sample tested at room temperature displayed cleavage planes associated with brittle fracture. However, the fracture surface of the sample tested at 300 °C (Fig. 5b) displayed cleavage planes as well as a large number of dimples which represented a mixture of brittle and ductile fracture at high temperature.

A-0-0 A-0-2 A-15-0 A-15-2 A-30-0 A-30-2

92 128 98 143 113 158

56 146 70 125 79 142

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Finally, the unique mechanical properties of CNTs and the special addition technique employed in this study have resulted in significant improvement of the mechanical properties as well as the fracture mode of A356-CNT composites. However, a number of studies have shown the possibility of formation of aluminum carbide (AI4C3) phase at the boundary zone of the CNTs and the matrix alloy [9, 20, 21]. Presence of interfacial phases can play an important role as a supporting medium to transfer stress from the matrix to the CNTs during the instability (necking) part of the deformation in tensile testing and can result in weakening of the composites. Study on the presence of AI4C3 phase at the boundary zone of the CNTs and the matrix alloy and the ways of preventing its formation will be the subject of our future works.

Figure 5. The fracture surface of composite sample A-30-2 after the tensile test at: (a) 25 °C and (b) 300 °C.

4. Conclusions A356 aluminum alloys reinforced with carbon nano-tubes (CNTs) were successfully produced by stir casting and a special compocasting routes. The employed compocasting route and injection of Ni-P-CNT coated aluminum particles instead of raw CNTs into the melt resulted in good distribution and less agglomeration of CNTs in the matrices of the produced A356-CNT composites which resulted in improved mechanical properties of the castings. Addition of CNTs to A356 matrix significantly improved such mechanical properties as Ys and UTS of the composite samples compared with those of the monolithic samples. While the UTS of the composite sample A-30-2 increased up to 195 MPa, its elongation decreased only 4% compared with that of the corresponding monolithic sample. Furthermore, the semi-solid composite samples kept more than 90% of their strength at 300°C, but the monolithic sample kept about 60% of its strength at this temperature.

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19. S. Tahamtan, M.A. Golozar, F. Karimzadeh, and B. Niroumand, "Microstructure and tensile properties of thixoformed A356 alloy," Materials Characterization, 59 (2008), 223-228.

20. H. Kwon, D.H. Park, J.F. Silvain, and A. Kawasaki, "Investigation of carbon nanotube reinforced aluminum matrix composite materials," Composites Science and Technology, 70 (2010), 546-550.

21. L. Ci, Z. Ryu, N.Y. Jin-Phillipp, and M. Rühle, "Investigation of the interfacial reaction between multi-walled carbon nanotubes and aluminum", Ada Mater, 54 (2006), 5367-5375.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

PRODUCTION OF CAST AZ91-CNT NANO-COMPOSITE BY ADDITION OF Ni-P-CNT COATED MAGNEIUSM POWDER TO THE MELT

Mahan Firoozbakht, Behzad Niroumand, Sayed Mahmoud Monirvaghefi

Department of Materials Engineering, Isfahan University of Technology, Isfahan 84156-83111, Iran

Keywords: Metal matrix composite, Carbon nanotube, AZ91, Nickel- phosphorus electroless coating on powder, stir casting, microstructure, mechanical properties.

Abstract

High agglomeration tendency and poor wettability of CNTs with metallic matrices are the main challenges for using this unique material as the reinforcement in magnesium based composites. In this study, nickel- phosphorous electroless coating process was employed for uniform co-deposition of CNTs into the Ni-P coating on the magnesium powder for the first time. The magnesium powder with the optimum composite coating was then used as the reinforcement and was introduced into the molten AZ91 alloy using stir casting method. Similar composites were also produced with untreated as-purchased CNTs for comparison. It was observed that addition of CNTs dispersed in the electroless coating resulted in refinement of the grain size, reduction in the porosity content and improvement of the hardness and flexural strength. Untreated CNTs were not able to improve the structural and mechanical characteristics of the cast composites to the same extent. This is believed to be due to the more agglomeration, less uniform distribution and poor wettability of the untreated CNTs with the melt compared to those introduced to the melt through the Mg-Ni-P-CNT composite powder.

Introduction

Magnesium and its alloys have been widely investigated because they have the lowest density among the industrial metals [1, 2] and have high specific mechanical properties [3, 4] which make them excellent choice for aerospace, automotive and sports applications [3-6]. Moreover, production costs for Mg alloys are comparatively low because of their good machinability, excellent cast ability and recognized recycling ability [2, 7], However, industrial usage of magnesium has some critical limitations. Elastic modulus and strength of magnesium is relatively low [1], In addition, strong directionality of properties due to hexagonal closed pack structure of magnesium limits its ductility and results in poor formability [3, 7], Nano-reinforcements have been shown to offer attractive solutions to such property shortcomings of magnesium [4, 5, 8]. Carbon nanotube (CNT) with its unique atomic structure and extraordinary mechanical properties is an ideal candidate as reinforcement for high strength, lightweight and high performance composites [9-12]. It has been reported that CNT can improve the strength and ductility of magnesium simultaneously [5, 13]. CNTs have large surface area (~1000nr/g) and much higher aspect ratio (~104) than traditional fillers for magnesium [9]. This high specific surface is desirable, in one hand, when it acts as an interface for an efficient stress transfer or when maximum resistance against dislocation movement is desired, but on the other hand, it

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causes strong tendency for CNTs to form aggregates due to the van der Waals interactions [9, 10, 14]. CNTs agglomeration is a great obstacle for their uniform distribution as reinforcement in composite materials [9, 13, 14]. The other important restriction is the limited wettability of the CNT with the metallic matrices [13-15]. Load transfer by the reinforcement depends on the perfect bonding at the interface and therefore a full properties improvement will not be accomplished due to this imperfection. To surmount these limitations, a two-steps process could be envisaged. In the first step, a homogenous dispersion of CNTs is achieved and then the second step for composite processing is carried out [9, 10, 14, 16]. For the first step, such methods as coating the CNTs with a more wettable metallic layer [1,16] or uniform dispersion of CNTs in a coating on metallic particles [9, 14, 17] have been suggested. By homogenous dispersion of CNTs into a metallic coating, such advantages as surface energy similarity of the metallic cover with that of the metal matrix and, subsequently, enhanced bonding at the interface can be achieved. The composite coating can be utilized in powder metallurgy [4, 6, 17] or liquid processing methods [5, 9, 10] to fabricate bulk metal matrix composites [18]. Ni-P electroless coating method have been used recently for producing metallic composite coatings by co-deposition of the reinforcement particles and the Ni-P matrix from a nickel electroless coating bath on a solid substrate [19,20]. Good dispersion and embedding of the CNTs co-deposited by this practice in the Ni-P matrix has been documented [21, 22], This co-deposition has the advantage of the metal coating around CNTs that improves the wettability and facilitates their uniform distribution as the reinforcement in the composites. Also, the Ni-P coating around CNTs will make their density closer to that of the lightweight alloy melts and so reduces their flotation to the surface during the processing. Accordingly, CNTs dispersion in a metallic coating can be a practical choice for the first step of nano-composites production. Subsequently and in the second step, a method for uniform introduction of the products of the first step to the melt should be employed. In a study carried out by the authors, electroless coating process was used for uniform co-deposition and dispersion of CNTs on Mg particles which were then introduced to molten AZ91 alloy by stir-casting process. The results of the electroless composite coating process have been published in [22]. The current paper presents the results of the second step of the process for production of cast AZ91-CNT nano-composites.

Experimental Procedure

The actual and standard composition of the matrix alloy used in this study is shown in Table 1.

Table 1. The Actual and Standard Composition of AZ91 Alloy Used as the Matrix Element, Wt% Actual Standard

Mg

89.77 Bal.

Al

9.27 8.1 -9.3

Mn

0.188 0.13

Zn

0.692 0.4-1.0

Si

0.0495 <0.30

Cu

<0.003 <0.10

Ni

<0.01 <0.01

Be

0.0003

Ca

0.0238 Other impurities

<0.30

Multi-walled carbon nanotubes produced by the Research Institute of Petroleum Industry (RIPI) in Tehran were used as the reinforcement. Commercially pure magnesium powder with mean diameter of 150 um was used as the coating substrate in the electroless coating process which

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employed Bath 18071 and Electroless Nickel SLOTONIP 70A process [23] from Schloetter Galvanotechnik Co. More details of the electroless composite coating process of Ni-P-carbon nanotubes on Mg powder and the optimum processing conditions have been stated elsewhere [22]. Stir casting process was then used for dispersion of the composite Mg powder into the AZ91 melt. The stir casting set-up used is schematically shown in the Figure 1. The melting and stir casting was carried out in an alumina crucible with bottom pouring arrangements fitted with steel baffles for eliminating the surface vortex. For preventing magnesium alloy from burning, a protection flux with the composition shown in Table II was used. The composite melt was finally bottom poured at 630°C in a steel mold placed under the furnace. The stirring speed in all the experiments was 500 rpm.

Figure 1. Schematic of the stir casting set-up used.

Table II. Composition of the Protection Flux Used Ingredients Wt%

MgCl2

40 CaF 20

CaCl2

14 MgO

12 NaCl

7 KC1

7

In different experiments, suitable amount of the composite Mg powder (Mg-Ni-P-CNT composite powder) to produce AZ91-CNT composites with 0.33, 0.67 and 1 percent carbon nanotubes were introduced to the AZ91 melt. For comparison, a monolithic sample was casted under exactly similar stir casting conditions. Furthermore, composite samples with similar CNT contents were made by introducing untreated as-purchased CNTs to the melt. Table III represents all the samples produced in this study. The cast samples were first degreased, cleaned and then their density was measured using the Archimedes principles as shown by Equation (1). Then, the porosity content of the samples was calculated using Equation (2). Microstructural characterization was conducted using a Nikon Model Epiphoto 300 optical microscope connected to an image analysis system. Nital 2% was used as etchant for metallographic observation of the samples. Grain size of the samples was measured using linear intercept method under ASTME112 standard.

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P = ms/((ms - mw) * pw) (1)

where p: Apparent density calculated from Archimedes principle ms: Mass of sample mw: Mass of sample floated in the water pw: Density of water (998.2 kg/m3 at 20°C)

P = (1- p/Ov * vr + pm . (1 - vr)) * 100 (2)

where P: Porosity (%) pr: Reinforcement density (PCNT= 13 gr/cm3) vr: Volume percent of reinforcement pm: Matrix density (PAZ91= 1.81 gr/cm3)

The average Vickers hardness as well as the flexural strength of the samples were measured. Flexural strength was measured by conducting three point bending test on round bars with diameter of 8.2 (±0.5) mm and using Equation 3. Error of the means [24] of the average grain size and hardness measurements reported in the corresponding tables was calculated using Equation 4.

oF=(8F*L)/(7T*D3) (3)

where OF: Flexural strength F: Ultimate load at three point bending test L: Length of specimen under load and between test apparatus legs D: Diameter of specimen

, = <r/vn (4)

where omean: Error of the mean a: Standard deviation n: Number of measurements

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Results and Discussions

1. Density and Porosity

Lightweight reinforcements such as CNT will reduce the density of the composites. However, due to the possible presence of porosity in the samples, the apparent density of the composites may be lower than their theoretical density. The calculated theoretical density (by the rule of mixtures) as well as the measured apparent density (by the Archimedes principle) and porosity content of the cast composites are shown in Table IV.

As the table clearly shows, the porosity content of the castings levels up by increasing the reinforcement content. This is due to the possible presence of a layer of air or other gases on the surfaces of CNTs as well as between the CNT agglomerates which is engulfed by the melt. It is logical to assume that such layers also exist on the surfaces of the composite Mg particles but in a lesser extent. This has been reflected in Table IV as the higher porosity content of the composites produced using un-treated CNTs compared to those produced using CNTs dispersed in the Ni-P coating. Also, higher agglomeration tendency of untreated CNTs is an important contributing factor to the higher porosity of the "P" samples.

2. Microstructure

Due to hexagonal close packed structure of magnesium, measuring its dendrites arm spacing and identifying the grains is not easy. So, the average linear intercept lengths of the microstructures are reported in Table V as a criterion of grain size.

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It is obvious from the results that the microstructures of the composites produced with introduction of composite Mg powder to the melt are finer than that of the monolithic sample and those produced using untreated CNTs. This shows that CNTs co-deposited into Ni-P electroless coating are more effective as grain refiner as the untreated CNTs. This can be due to the better distribution and less agglomeration of the CNTs as well as the improved wettability of the Ni-P coated CNTs and the reinforcement/melt contact. This is expected to results in better retaining of the CNTs in the matrix. As the table shows, one percent of CNTs introduced in the form of Ni-P-CNT coating has resulted in 27% reduction in the average linear intercept length of the grains, while one percent of pure CNTs resulted in only 13% reduction in the grain size.

3. Hardness

The average measured hardness of the cast samples are presented in Table VI. As the table shows, the hardness has not changed significantly by addition of 0.33% CNT in any form. However, it increases by increasing the CNT content of the samples in both conditions. It's clear that CNTs have been more effective for improving the hardness when introduced into the melt through the composite Mg powder. One percent of CNTs dispersed in the coating has increased the hardness by about 36 percent, while one percent of untreated CNTs has only improved the hardness by about 11 percent. This is believed to be mainly due to the better wetting of the reinforcement with the matrix, more uniform distribution of the reinforcement and the higher retainment of the CNTs in the matrix.

4. Flexural Strength

The average flexural strengths of the cast samples are shown in Table VII.

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It is clear that the flexural strength increases by increasing the CNT content of the samples. Also, CNTs have been more effective in improving the strength when introduced into the melt through the composite Mg powder. One percent of CNTs dispersed in the coating has improved the flexural strength by about 30 percent, while one percent of untreated CNTs has improved the strength by 21 percent.

Conclusions

Production and characterization of CNT reinforced AZ91 a magnesium alloy composite was investigated in this work. For this purpose, a two-step process was employed. The first step included the co-deposition of CNTs with Ni-P electroless coating on magnesium powder. The Ni-P-CNT coated magnesium powder was then introduced into the melt by stir casting method and resulted in refinement of the grain size, reduction in the porosity content and improvement of the hardness and flexural strength. Untreated CNTs were also added to the melt under the similar conditions, however were not able to improve the structural and mechanical characteristics of the cast composites to the same extent. This is believed to be due to the more agglomeration, less uniform distribution and poor wettability of the untreated CNTs with the melt compared to those introduced to the melt through the composite magnesium powder.

References

1. T. Honma, K. Nagai, A. Katou, K. Arai, M. Suganuma, S. Kamado, "Synthesis of high-strength magnesium alloy composites reinforced with Si-coated carbon nanofibers", Scripta Materialia, 60 (2009), 451-454. 2. Z.H. Xiao, J.R.Luo, S.S.Wu, D.N. Li, Y.W. Mao, X.J. Song, "Study on a semi-solid rheo-diecasting process with AZ91D alloy slurry", Journal of Materials Engineering and Performance, 13 (2004), 60-63 3. C.S. Goh, J. Wei, L.C. Lee, M. Gupta, "Ductility improvement and fatigue studies in Mg-CNT nanocomposites", Composites Science and Technology, 68 (2008), 1432-1439. 4. S.K. Thakur, T.S. Srivatsan, M. Gupta, "Synthesis and mechanical behavior of carbon nanotube-magnesium composites hybridized with nanoparticles of alumina", Materials Science and Engineering A, 466 (2007), 32-37. 5. C.S. Goh, J. Wei, L.C.Lee, M. Gupta, "Simultaneous enhancement in strength and ductility by reinforceing magnesium with carbon nanotubes", Materials Science and Engineering A, 423 (2006), 153-156. 6. C.S. Goh, J. Wei, L.C. Lee, M. Gupta, "Development of novel carbon nanotubes reinforced magnesium nanocomposites using powder metallurgy techniques", SIMTech technical reports, 9 (2008), 130-135. 7. S.F. Hassan, M. Gupta, "Development of ductile magnesium composite materials using titanium as reinforcement", Journal of Alloys and Compounds, 345 (2002), 246-251. 8. M.A. Thein, L. Lu, M.O. Lai, Effect of milling and reinforcement on mechanical properties of nanostructured magnesium composite", Journal of Materials Processing Technology, 209 (2009), 4439-4443. 9. Q. Li, A. Viereckl, CA. Rottmair, R.F. Singer, "Improved processing of carbon nanotube/magnesium alloy compsoites", Composites Sience and Technology, 69 (2009), 1193-1199.

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10. X. Zeng, G. Zhou, Q. Xu, Y.Xiong, C.Luo, J. Wu, "A new technique for dispersion of carbon nanotubes in a metal melt", Materials Science and Engineering A, 527 (2010) 5335-5340. 11. A. Esawi, K. Morsi, "Dispersion of carbon nanotubes (CNTs) in aluminum powder", Composites: Part A, 38 (2007), 646- 650. 12. 1. Bright, V. Koutsos, Q. Li, R. Cheung, "Carbon nanotubes for integration into nanocomposite materials", Microelectronic Engineering 83 (2006), 1542- 1546. 13. L. Shi-ying, G. Fei-peng, Z. Qiong-yuan, Z. Xue, L. Wen-Zhent, "Fabrication of carbon nanotubes reinforced AZ91D composites by ultrasonic processing", Transactions ofNonferrous Metals Society of China, 20 (2010), 1222-1227. 14. K. Kondoh, H. Fukuda, J. Umeda, H. Imai, B. Fugetsu, M. Endo, "Microstructural and mechanical analysis of carbon nanotube reinforced magnesium alloy powder composites", Materials Science and Engineering A, 527 (2010), 4103-4108. 15. H. Uozumi, K. Kobayashi, K. Nakanishi, T. Matsunaga, K.. Shinozki, H. Sakamoto, T. Tsukada, C. Masuda, M. Yoshida, "Fabrication process of carbon nanotube/light metal matrix composites by squeeze casting", Materials Science and Engineering A, 495 (2005) 282- 287. 16. E. Carreflo-Morelli, J. Yang, E. Couteau, K. Hernadi, J.W. Seo, C. Bonjour, L. Forrô, R. Schaller, "Tiny tubes boost for metal matrix composites", Physica Status Solidi (a), 201 (2004), R53- R55. 17. Y. Shimizu, S. Miki, T. Soga, I. Itoh, H. Todoroki, T. Hosono, K. Sakaki, T. Hayashi, Y.A. Kim, M. Endo, S. Morimoto, A. Koide, "Multi-walled carbon nanotube-reinforced magnesium alloy composites", Scripta Materialia, 58 (2008), 267-270. 18. B. Abbasipour, B. Niroumand, S.M. Monir Vaghefi, "Compocasting of A356-CNT composite", Transaction Nonferrous Metals Society of China, 20 (2010), 1561-1566. 19. Z. Yang, H. Xu, M.K. Li, Y.L. Shi, Y. Huang, H.L. Li, "Preparation and properties of Ni/P/single-walled carbon nanotubes composite coatings by means of electroless plating", Thin Solid Films, 466 (2004), 86-91. 20. C.S. Chen, X.H. Chen, Z. Yang, W.H. Li, L.S. Xu, B. Yi, "Effect of multi-walled carbon nanotubes as reinforcement fibers on tribological behavior of Ni-P electroless coating", Diamond Related Materials, 15 (2006), 151-156. 21. Y.T. Wu, L. Lei, B. Shen, W.B. Hu, "Investigation in electroless Ni-P-Cg(graphite)-SiC composite coating", Surface Coating Technology, 201 (2006), 441-445. 22. M. Firoozbakht, S.M. Monirvaghefi, B. Niroumand, "Electroless composite coating of Ni-P-Carbon nanotubes on magnesium powder", Journal of Alloys and Compounds, 509S (2011), S496- S502. 23. 18071 Electroless Nickel SLOTONIP 70 A.pdf, www.schloetter.de/en/processes/18-electroless-nickel/electroless-nickel/ (Last visited on 26.06.2010) 24. J. Topping, Errors of Observation and their Treatment, 4th ed. (Chapman & Hall, 1972).

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Supplemental Proceedings: Volume 1 : Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

WEAR BEHAVIOR OF MAGNESIUM MATRIX NANOCOMPOSITES AT ROOM AND ELEVATED TEMPERATURE

Wenzhen Li, Shiying Liu

Department of Mechanical Engineering, Tsinghua University, Beijing 100084, China

Keywords: magnesium matrix nanocomposites, dry sliding wear, AZ91D alloy, SiC nanoparticles

Abstract

Magnesium matrix nanocomposite reinforced with nano-sized SiC particles was fabricated by mechanical stirring and high intensity ultrasonic dispersion processing. The dry sliding wear behavior of the magnesium alloy and its nanocomposites were investigated using a ball-on-disc CETR wear testing machine. The test results show that the addition of SiC nanoparticles increased the wear resistance of composites. The wear loss increase as the load increasing as expected and the wear loss of AZ91D is higher than those of its composites at all the loads. The dominative wear mechanisms vary with different load and the load occurring delamination wear can be improved due to SiC nanoparticles strengthening function. Wear test results also showed that the worn stages of AZ91D and its composite are: running and stable worn stages without occurring acutely worn stages in one hour sliding time.

Introduction

As the energy and environmental problem more and more serious in the world, the development trend of the industry is to save energy and the green production. Among all the metal engineering materials, the research and development of iron, aluminum and copper are getting mature, which cause their resources reserves is declining. It is necessary to prospect new resource. Magnesium and its alloys, as one of the lightest industrial metals in the practical using, have become a great deal of alternate resources. However, their mechanical and friction properties are poor. In particular, when applying them to friction materials, the wear or seizure phenomena easily occur by contacting with the counter materials [1]. In order to cover these deficiencies of magnesium, efforts have been made to develop Mg metal matrix composites. Magnesium matrix composite is a kind of excellent candidate material for the applications of aerospace applications, automotive and civilian fields for the advantages of high specific strength, specific stiffness, good wear resistance and low density. Using nano-sized particles as reinforcements to fabricate metal matrix composites, a small volume fraction of nanoparticles can significantly improve the reinforced effect [2], Many works have been done on fabrication and mechanical properties of magnesium matrix nanocomposites [3-6] , while less work has been carried out on the wear behavior of nanocomposites.

Many studies have been done on the wear of magnesium composites reinforced with micro-sized particles, fibers, CNTs or hybrid component [7-12]. The wear behavior is affected on the shapes, sizes and fraction volume of the reinforcements .The reinforcements distribution, interface between the reinforcement and matrix also affects the tribological behavior of composites. The testing condition, such as the type of coûter surface, normal load, sliding speed and temperature also affects the tribological behavior of composites. CNTs having good self-lubricating property and a friction coefficient between the walls of multi-walled CNTs should be extremely low [10].

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CNTs had an important role to form the lubricant condition at the sliding interface between the matrix and the counter material due to their network structure on the sliding surfacefl 1]. The wear resistance of composites increases with the decrease of the particle size [12] The nanocomposites processes good wear resistance [13,14], Lim et al [13]studied wear of magnesium composites reinforced with nano-sized alumina particulates, they found that the wear resistance of the composites improved with increasing amounts of reinforcement. Only l.llvol% of nano-sized alumina particulates was effective up to 1.8 times increase in the wear resistance of pure magnesium. Habibnejad-Korayem M etal [14] studied dry sliding wear behavior of Mg and AZ31 strengthened by A1203 nano-particles. Although grain refinement, higher load-bearing capacity, and improved hardness have some contributions to the overall wear resistance, increased work-hardening capacity due to the interaction of dislocations and nano-particles can be considered as the main mechanism improving wear behavior of the present nano-composites.

With such limited studies conducted, these findings suggest that the composites reinforced with nanoparticles processes good wear resistance. In this study, magnesium matrix composites reinforced by dispersed SiC nanoparticles are fabricated by mechanical stirring and high-intensity ultrasonic dispersion processing. The dry sliding wear behavior of the magnesium alloy and its nanocomposites were investigated at room temperature and at elevated temperature.

Experiment

Materials and Fabrication Procedure

The magnesium matrix composite with different weight percent nano-sized SiC particles were fabricated by mechanical stirring and high-intensity ultrasonic dispersion method. The matrix material was Mg-Al-Zn alloy (AZ91D). The reinforcements were SiC particles with average diameter of 40 nm.

Steps involved and the procedure employed for the fabrication of composites are as follows: the melting was carried out using a resistance heating furnace. A clay-graphite crucible with 130mm in inner diameter and 180mm in height was used to melt magnesium alloy. One time about 1.8 kg AZ91D was melted in clay-graphite crucible. In order to avoid oxidation and burning, the magnesium melt was protected by CO2+0.2vol.%SF6. When the alloy was completely melted, the mechanical stirring was applied and SiCp were added into the stirring vortex. The stirring speed was 300-500 r/min and stirring time was about 1-2 min. In order to disperse better SiCp within the matrix, a high-intensity ultrasonic wave with a 15 kHz, a maximum 2.6 kW power output and a titanium alloy waveguide of 40 mm in diameter was used for processing SiCp/AZ91D melt. The ultrasonic processing time was 15 min and the ultrasonic probe was dipped into the melt for about 30 mm. After ultrasonic processing, the composite melt was heated to 660 'C then cast into a steel permanent mold (preheated to 250-300 °C). For comparison, samples without the addition of SiCp were prepared.

Wear tests

Wear tests were carried out in dry sliding condition using a ball-on-disc CETR UMT 3-V wear testing machine. The counterbodies were 440C stainless steel with 9.5 mm in diameter and hardness of 62HRC. The disc wear test specimens of 50 mm in diameter and 9 mm in height were machined from castings. The wear tests were carried out at the lod of 10, 30, 50 and 70N respectively. The disc rotates horizontally at sliding speed of 0.2m/s and the sliding time was one

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hour. The tests were made at room temperature and at elevated temperature up to 300 °C. Weight loss was measured by using an electronic balance with a resolution of 0.01 mg. Before and after each test, the disc specimens were cleaned in an ultrasonic bath with acetone and were dried in the air. The worn surfaces were investigated with a scanning electron microscopy (FEI Siron200 SEM).

Results and Discussion

Microstructures and Mechanical Properties

The dispersion of SiC nanoparticles in the magnesium matrix was researched by SEM, as shown in Figure 1. The obtained magnesium matrix nanocomposites are uniform and dense in microstructure without obvious shrinkage porosity. Microstructural studies conducted on the composites specimens shown uniform reinforcement distribution. It is clear that SiC nanoparticles were dispersed and distributed very well throughout the magnesium matrix. The well-dispersed SiC nanoparticles can improve the mechanical properties of composites.

Table I shows that the strength and hardness of nanocomposites increases with the increase in weight percentage of SiC nanoparticles. The density of nanocomposites decreases with increasing weight percentages of SiC and the density in all the nanocomposites is lower than that in AZ91D. But the density of all nanocomposites is higher than 98%. Adding SiC nanoparticles, the grain size gradually decreases and the average grain size is reduced by 71.9% for 2.0 wt.%SiCp/AZ91D.

Figure 1. SEM images of AZ91D based nanocomposites

Table I. Properties of AZ91D and Its Nanocomposites Mateials UTS/MPa Hardness/HV Density/% Grain size/um AZ91D 107 66.6 99.72 562.32 0.5wt.%Sicp/AZ91D 114 70.6 99.45 327.51 1.0wt.%Sicp/AZ91D 136 73.1 99.01 283.57 1.5wt.%Sicp/AZ91D 137 74.4 98.96 175.86 2.0wt.%Sicp/AZ91D 142 77J! 98.41 157.87

The Effects of SiC Nanoparticles on Wear Behavior

Figure 2 shows the worn surfaces of AZ91D and its composites at sliding speed of 0.2 m/s at 50 N load for sliding time up to one hour. The worn surface of AZ91D reveals obvious plastic deformation and continuous wide and deep groove interspersed by parallel small grooves along sliding direction. This is typical abrasive wear, in which hard particles or hard asperities plough or cut into friction pair and form scratches along the sliding direction. From Figure 1 (a), we also found that there exist adhesive wear and delamination wear. It shows the severe wear regime for

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AZ91D matrix. Figure 1 (b-d) shows that the worn surface of composites with different weight percent nano-sized SiC particles. The smooth worn surfaces of composites contains a few shallow grooves. One explanation is that SiC nanoparticles are peeled from matrix during wear and form three body abrasion with debris, as shown in Figure 3. The wear resistance of nanocomposites is higher than that of pure AZ91D alloy. This is mainly due to that (1) the grain refinement strengthening and the Orowan strengthening owing to the exist of the SiCp [15]; (2) as the coefficient of thermal expansion and the distortion of the SiCp differs greatly to the AZ91D alloy, the dislocation density will be increased when the plastic deformation happens in the process of friction, thus the work-hardening capacity is increased [14].

Figure 2. SEM images of worn surface of AZ91D and its nanocomposites

Figure 3. Abrasive on worn surface of Figure 4. The effect of load on wear loss 1 0"/ SiCü/AZ91D of AZ9ID and its nanocomposites

The Effects of Load on Wear Behavior

The wear loss for AZ91D and its composites are plotted against the load in Figure 4. The results show that the wear loss increase as the load increasing as expected. The 0.5wt.%SiCp/AZ91D shows an increasing in the wear loss of 3 times at 70 N load than that at ION load. It is also apparent that the wear loss of AZ91D is higher than those of its composites at all the loads. The wear loss of AZ91D is 27.3% higher than 2.0wt.%SiCp/AZ91D composite at 10 N load and 10.6% higher at 70 N load. The wear rate is increased with the sliding distance and load according to Archard equation [16]. Under the same sliding speed, the greater the load is, the heavier the wearing is.

The worn surfaces of AZ91D and 1.0wt.%SiCp/AZ91D at different loads are shown in Figure 5 and Figure 6 respectively. The worn surface of AZ91D exhibits plastic stacks and adhesion ( Figure5a), and that of 1.0wt.%SiCp/AZ91D are adhesion and discontinuous small grooves ( Figure 6a). At the lower loads, the adhesive wear was the dominative wear mechanism. With the increase of the load, the worn surfaces of materials reveal obvious plastic deformation and continuous grooves. At 30 N load, the worn surface of AZ91D exists fatigue cracks and more

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broad grooves than that of 1.0wt.%SiCp/AZ91D, the abrasive wear was the dominative wear mechanism. There exist more broad grooves in the worn surface of AZ91D at 50N, severely wear occurs and then coating scaling presents in lamellar shape by delamination mechanism. There were also many parallel grooves and small cracks in the worn surface of 1.0wt.%SiCp/AZ91D, the delamination wear will occur and the abrasive wear was the dominative wear mechanism. The load occurring delamination wear can be improved due to SiC nanoparticles strengthening function. When the load increases to 70 N, the delamination wear occur for 1.0wt.%SiCp/AZ91D and the more seriously delamination wear occur for AZ91D.

Figure 5. SEM micrograghs of worn surface of AZ91D at different loads

Figure 6. SEM micrograghs of worn surface of 1.0wt.%SiC/AZ91D at different loads

The Effects of Sliding Time on Wear Behavior

Figure 7and Figure 8 shows the effects of sliding time on coefficient of friction. The worn stages of AZ91D and its composite are: running and stable worn stages without occurring acutely worn stages in one hour sliding time. The coefficient of friction for composite is lower than that of the matrix due to addition nano-sized SiC reinforcement.

Figure 7 shows the effects of sliding time on coefficient of friction at 50N load and speed of 0.4m/s. It depicts that friction coefficient of sample tended to reduce and stabilize with the extending of slidmg time. The coefficient of friction decreases obviously with increase in sliding time in the initial stage, which is running stage. It takes 380s enter stable worn stage and the coefficient of friction for AZ91D was stabilized at about 0.28. The coefficient of friction for composite was stabilized at about 0.26 after 140s running worn. When the sliding speed decreased to 0.2m/s from 0.4m/s, the coefficients of friction for AZ91D and 0.5SiC/AZ91D decrease as the sliding time increases and enter stable worn stage after long time running (Figure 8) The coefficients of friction of stable worn stage for AZ91D and 0.5SiC/AZ91D at

speed of 0.2m/s was about the same as that of speed of 0.4m/s.

Wear firstly arises on the asperities of two metal surfaces. SiC reinforcements due to their nanometer size can not bear load to reduce the force of composites. So in initial wear stage, the

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asperities of materials contact the surface of counterbodies, in which the coefficient of friction is in a gradually decreasing process.. The worn surface occurred repeatedly plastic deformations and collapse, the working hardening gradually increases and becomes stable with the increase of the worn time [14]. When the working hardening reaches a certain element to the worn surface covered with adhesive transference, the worn stages would change from running to stable worn. In stable worn stage, the coefficient of friction only slightly affects by sliding time. The friction coefficient of nanocomposite is lower than that of AZ91D in whole worn stage. This is because namoparticles can decrease the adhesion between the couple parts, which rapidly enters in stable wear stage. Moreover, the coefficient of friction can be decreased due to nanoparticles strengthening function and tight interface between nanoparticles and matrix.

Figure 7. The effects of sliding time on coefficient of friction at 50N load with sliding speed

of0.4m/s

Figure 8. The effects of sliding time on coefficient of friction at 50N load with sliding speed of 0.2m/s

The Effects of Temperature on Wear Behavior

Figure 9 shows the effects of wear temperature on coefficient of friction. Figure 10 shows the effects of wear temperature on weight loss. The wear resistance of nanocomposites is better than that of pure AZ91D alloy due to nanoparticles reinforced. The weight loss of the nanocomposites sample increases with increase in load for the matrix and its nanocomposites as the rising temperature. Temperature has important influence on the wear-resisting property of materials. From room temperature to 300°C, the AZ91D and its composites undergo mild-to-severe wear

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regime at different temperature. The transition temperature of the nanocomposites from mild-to-sever wear increases 50°C than that of AZ91D matrix (Figure 10). The nanocomposites show better resistance to high-temperature wear property.

Figure 9 The effects of wear temperature on Figure 10 The effects of wear temperature on coefficient of friction weight loss

Conclusion

The SiC/AZ91D nanocomposites with different weight fraction of reinforcing SiC nanoparticles were fabricated by mechanical stirring and high-intensity ultrasonic dispersion processing. The dry sliding wear tests with AZ91D and its composites were carried out at different load. The following conclusion can be draw from the experimental results:

• The wear loss is increased with the load from ION to 70N. The wear resistance of the alloy is improved significantly due to nanoparticles addition. The wear loss of AZ91D is higher than those of its composites at all the loads.

• The worn mechanism of AZ91D and its composites are adhesion wear with slightly abrasive wear at low load of ION. With the increasing of the test load, the some wear debris as abrasive cut the surface form grooves, the wear mechanism changes from adhesion wear to abrasive wear (30N load). When load was further increased to 50N, the severely wear occurs and then coating scaling presents in lamellar shape by delamination mechanism for AZ91D and abrasive wear was the dominative wear mechanism for nanocomposites. When the load increases to 70 N, the delamination wear occur for 1.0wt.%SiCp/AZ91D and the more seriously delamination wear occur forAZ91D.

• The worn stages of AZ91D and its composite are: running and stable worn stages without occurring acutely worn stages in one hour sliding time.

• The improvement of the wear resistance of nanocomposites is mainly due to the grain refinement strengthening and the Orowan strengthening. Moreover, as the coefficient of thermal expansion and the distortion of the SiCp differs greatly to the AZ91D alloy, the dislocation density will be increased when the plastic deformation happens in the process of friction, thus the work-hardening capacity is increased

• The nano particles can improve the wear resistance of nanocomposites. The transition temperature of the nanocomposites from mild-to-sever wear increases 50°C than that of AZ91D matrix . The nanocomposites show better resistance to high-temperature wear property.

References

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1. N.N. Aung, W. Zhou, L. E.N. Lim, Wear behaviour of AZ91D alloy at low sliding speeds, Wear, 265 (2008), 780-786. 2. G.Cao, H.Konishi, X.Li, Mechanical properties and microstructure of SiC-reinforced Mg-(2,4)Al-lSi nanocomposites fabricated by ultrasonic cavitation based solidification processing. Materials Science and Engineering ^,486(2008),357-362. 3. C.S.Goh, J.Wei, L.C.Lee, M.Gupta. Development of novel carbon nanotube reinforced magnesium nanocomposites using the powder metallurgy technique, Nanotechnology, 17(2006), 7-12. 4. L. Lu, M.O. Lai, W. Liang, Magnesium nanocomposite via mechanochemical milling, Composites Science and Technology, 64 (2004), 2009-2014. 5. H. Ferkel, B.L. Mordike, Magnesium strengthened by SiC nanoparticles, Materials Science and Engineering A, 298 (2001) 193-199. 6. J.Lan, Y.Yang, X. C. Li, Microstructure and microhardness of SiC nanoparticles reinforced magnesium composites fabricated by ultrasonic method, Materials Science and Engineering A, 386(2004), 284-290. 7. A. Alahelisten, F. Bergman, M. Olsson, S. Hogmark, On the wear of aluminum and magnesium metal matrix composites, Wear, 165 (1993) 221-226. 8. S.C. Sharma, B. Anand, M. Krishna, Evaluation of sliding wear behaviour of feldspar particle-reinforced magnesium alloy composites,Wear, 241 (2000) 33-40. 9. S.K. Thakur, B.K. Dhindaw, Influence of interfacial characteristics between SiCp and Mg/Al metal matrix on wear, coefficient of friction and microhardness, Wear, 247 (2001), 191-201.

10. T. K. Kyung, I. C. Seung, H.H. Seung, Microstructures and tensile behavior of carbon nanotube reinforced Cu matrix nanocomposites, Materials Science and Engineering A, 430 (2006) 27-33. l l . J . Umeda, K. Kondoh, H. Imai, Friction and wear behavior of sintered magnesium composite reinforced with CNT-Mg2Si/MgO, Materials Science and Engineering A, 504 (2009), 157-162. 12. M. Aydin, F. Findik, Wear properties of magnesium matrix composites reinforced with SiCh particles, Industrial Lubrication and Tribology, 62 (2010), 232-237.

13. C.Y.H. Lim, D.K. Leo, J.J.S.Ang, A.M.Gupta, Wear of magnesium composites reinforced with nano-sized alumina particulates. Wear, 259(2005) 620-625. 14. M. Habibnejad-Korayem, R.Mahmudi, H.M.Ghasemi, W.J. Poole, Tribological behavior of pure Mg and AZ31 magnesium alloy strengthened by AI2O3 nano-particles. Wear 268(2010),405-412. 15. Z.Zhang, D.L. Chen, Consideration of Orowan strengthening effect in particulate-reinforced metal matrix nanocomposites: A model for predicting their yield strength. Scripta Materialia, 54(2006), 1321-1326. 16. J.F. Archard. Contact and rubbing of flat surfaces. Journal of Applied Physics, 24 (1953), 981-988.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals <£ Materials Society), 2012

Uniform Dispersion of Nanoparticles in Metal Matrix Nanocomposites

L.Y. Chen, H. Choi, A. Fehrenbacher, J.Q. Xu, C. Ma, X.C. Li

Department of Mechanical Engineering, University of Wisconsin-Madison; 1513 University Avenue; Madison, WI 53706, USA

Keywords: Metal matrix nanocomposites, Aluminum alloy, Ultrasonic processing, Friction stir processing

Abstract

The dispersion of nanoparticles in metal matrices is crucial for the performance of the metal matrix nanocomposites (MMNCs). Uniformly dispersing nanoparticles into the metal matrix is still a great challenge. This paper demonstrates that a uniform dispersion of AI2O3 nanoparticles in an A206 alloy matrix is achieved by a combination of liquid state ultrasonic processing and solid state friction stir processing. The results obtained in this work provide a general methodology to uniformly disperse nanoparticles into metal matrix nanocomposites, paving the way for development of metal matrix nanocomposites with superior properties.

Introduction

Metal matrix composites (MMCs), usually a combination of continuous metallic phase and ceramic particles/fibers, have potential to possess the merits of both metals and ceramics [1-7]. Traditionally, micro-sized particles are used to fabricate MMCs [1-3, 7]. The micro-sized particle reinforced MMCs exhibit higher strength than their base alloys without particles, gaining the advantages of ceramics; unfortunately, the micro-sized particle reinforced metal matrix composites also inherit the brittleness of ceramics, exhibiting poor ductility [8]. Recent studies show that materials become insensitive to flaws at the nanoscale [9]. Uniformly embedding nano-sized ceramic particles into the metal matrix to fabricate metal matrix nanocomposites (MMNCs) might achieve dramatic enhancement of strength without loss of ductility, realizing the full potential of metal matrix composites [5, 10]. During the last decade, considerable effort has been directed to the development of metal matrix nanocomposites. Several methods (powder metallurgy [11], vortex process [12], accumulative roll bonding [13], etc.) are developed to introduce nanoparticles into the metal matrix. However, uniformly dispersing nanoparticles into the metal matrix is still a huge challenge.

Recently, ultrasonic processing is being developed to incorporate and disperse nanoparticles into molten metals [14, 15]. It is believed that the transient cavitation (with temperatures of about 5000 °C, pressures of about 1000 atm) and acoustic streaming can incorporate and disperse the nanoparticles in the liquid state effectively [14, 15]. However, during solidification, the dispersed nanoparticles, especially alumina, in the liquid state, are often pushed to the grain boundaries due to the repulsive Van der Waals force between the solidification front and nanoparticles [16]. Thus, in the solidified state, the nanoparticles tend to distribute along grain boundaries. This is a great challenge for all liquid state-based processing methods.

In 1991, a solid-state joining technique - friction stir welding (FSW) - is invented at The Welding Institute (TWI, Cambridge, United Kingdom) [17], The joining is accomplished by inserting a nonconsumable rotating tool with a specially designed pin and shoulder into the abutting edges of plates. The tool travels along the joint line, heats the workpieces, and moves the materials around the pin to produce the joint in solid state. Recent studies show that this

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technique can mix the materials very well in solid state and modify the microstructure (friction stir processing, FSP) [18-20]. Several groups have tried to prepare metal matrix nanocomposites using the mixing effect of this technique. The particles are usually introduced by spraying them on the surface or placing them into a grove. Even though multiple passes are used to disperse the particles, the particle distribution is not uniform at a large area. Moreover, the nature of the nanoparticle incorporation during friction stir processing limits its potential for large scale production of MMNCs.

In this work, we experimented to develop a novel method combining the advantages of both liquid state ultrasonic processing and solid state friction stir processing to disperse ceramic particles uniformly in the metal matrix.

Figure 1 Schematic of the processing procedure to obtain uniform dispersion of nanoparticles in the metal matrix.

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Aluminum alloy A206 is selected as the matrix. Alpha AI2O3 nanoparticles with a size of about 100-200 nm are selected as reinforcement particles. To disperse the reinforcement particles uniformly into the A206 matrix, a novel method combining liquid state ultrasonic processing and solid state friction stir processing is developed. The schematic of the method is shown in Figure 1. First, AI2O3 nanoparticles are fed into the melt of A206 with a double-capsulate feeding method [21], and are then dispersed by ultrasonic processing. Figure la shows the schematic of the setup for liquid state ultrasonic processing. The system consists of a resistance heating furnace, a protection gas system (argon), a nanoparticle feeding system, and an ultrasonic processing system, which is composed of an ultrasonic probe made of niobium alloy C103, a booster, and a transducer. After ultrasonic processing, it is expected that the particles are uniformly dispersed in the melt. Then the melt with uniformly dispersed AI2O3 particles is cast to a plate mold. After solidification, the AI2O3 particle reinforced A206 plate with thickness of 6 mm is obtained, as shown in figure lb. However, the particle distribution in the plate is not uniform. Due to the repulsive Van der Waals force between the AI2O3 particles and the solidification front, the AI2O3 particles are pushed to the grain boundaries [16],

Figure 2 SEM images of the as-cast ultrasonic processed plate (a, b) and the ultrasonic processed and solid state stirred sample (c, d). The white spots in the images are AI2O3 particles.

To distribute the particles uniformly in the matrix, the as-cast plate is stirred in solid state by friction stir processing. Figure lc demonstrates the details of this solid state stirring. A rotating pin with diameter of 5 mm and length of about the thickness of the plate is inserted into the plate and travels across the plate. The pin with rotating speed of 900 rpm stirs the materials along the travelling path with a travel speed of 25 mm/min in solid state, as shown in figure lc. The dispersion of the particles in the stirred region is observed by scanning electron microscope (SEM). The tensile bars with thickness of 3.5 mm, width of 5 mm and gauge length of 31.75 mm

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are cut from the stirred region (processed region parallel to tensile direction) and tested by a mechanical testing machine (Sintech 10/GL, MTS) with a crosshead speed of 5.08 mm/min.

Results and Discussion

Figure 2a shows the distribution of the AI2O3 particles in the as-cast plate. It is found that the distribution of particles is uniform but the dispersion is not satisfactory. The particles tend to form a network in the grain boundaries. Previous study indicates that this phenomenon results from the repulsive interaction between AI2O3 particles and the solidification front [16]. Figure 2b is the high magnification image of the high density particle region. Even though some micro clusters still exist, most of the particles are dispersed very well. This implies that the particles are dispersed very well in liquid state by ultrasonic processing. However, the well dispersed particles in liquid state are pushed together by solidification front during solidification due to the repulsive Van der Waals force between the AI2O3 particles and the solidification front.

Figure 3 Engineering stress-strain curves of the pure A206 alloy, the ultrasonic processed nanocomposite, and the ultrasonic processed and solid state stirred nanocomposite.

To disperse the particles further, solid state friction stir processing is conducted. Figure 2c demonstrates the particle distribution in the stirred region. It can be seenthat the particles distribute uniformly in the specimen. The high magnification image in figure 2d confirms that the particles are well dispersed and distributed in the stirred region.

To check the mechanical properties of the stirred region with uniformly dispersed AI2O3 particles, tensile bars with a thickness of 3.5 mm, a width of 5 mm and a gauge length of 31.75 mm are cut from the stirred region by electrical discharge machining. The tensile bars are tested by a tensile testing machine with a crosshead speed of 5.08 mm/min. For comparison, the as-cast ultrasonic processed nanocomposite and pure A206 alloy are also tested. The engineering stress-strain curves are shown in figure 3. The yield strength, ultimate tensile strength, and elongation of the tested samples are listed in table 1. The results show that the solid stirred sample with uniform dispersion of AI2O3 nanoparticles exhibits both enhanced strength and ductility. Further study is needed to unravel the mechanism of the dramatic enhancement of strength and ductility induced by solid state stirring.

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Table 1 Yield strength, tensile strength, and elongation of the A206 alloy and their nanocomposites

Sample Yield strenth Tensile strength Elongation

(MPa) (MPa) (%) A206as-ast 111 173 3.3 A206 + A1203 ultrasonic 119 221 7.3 A206 + AI2O3 ultrasonic + FSP 180 315 18

In order to further confirm the uniform dispersion of the particles after solid state stirring, we observed the fracture surface of the stirred sample using SEM. The images demonstrate that the particles distribute uniformly on the fracture surface, as shown in figure 4.

Figure 4 SEM image of the fracture surface of the ultrasonic processed and solid state stirred nanocomposite.

Conclusions

A method integrating liquid state ultrasonic processing and solid state friction stir processing is developed to incorporate, distribute and disperse AI2O3 nanoparticles into an A206 alloy matrix. The results show that uniform dispersion of AI2O3 nanoparticles in the A206 alloy matrix is achieved. The ultrasonic processed and solid state stirred nanocomposites exhibit both enhanced strength and ductility. The results obtained in this work provide a general methodology to uniformly disperse nanoparticles into metal matrix nanocomposites, paving the way for development of metal matrix nanocomposites with enhanced properties.

Acknowledgements

This work is supported by the National Institute of Standards and Technology through its Technology Innovation Program. Help with friction stir processing experiments from Professor Frank E. Pfefferkorn and his lab at the University of Wisconsin-Madison is greatly appreciated.

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References

1. I.A. Ibrahim, F.A. Mohamed, and E.J. Lavernia, "Paniculate Reinforced Metal Matrix Composites - A Review," Journal of Materials Science, 26 (1991), 1137-1156.

2. A. Kelly, "Composites in Context," Composites Science and Technology, 23 (1985), 171-199.

3. A. Mortensen and I. Jin, "Solidification Processing of Metal Matrix Composites," International Materials Reviews, 37 (1992), 101-128.

4. A. Mortensen and J. Llorca, "Metal Matrix Composites," Annual Review of Materials Research, 40 (2010), 243-270.

5. S.C. Tjong, "Novel nanoparticle-reinforced metal matrix composites with enhanced mechanical properties," Advanced Engineering Materials, 9 (2007), 639-652.

6. S.C. Tjong and Z.Y. Ma, "Microstructural and mechanical characteristics of in situ metal matrix composites," Materials Science & Engineering R-Reports, 29 (2000), 49-113.

7. D.J. Lloyd, "Particle-Reinforced Aluminum and Magnesium Matrix Composites," International Materials Reviews, 39 (1994), 1-23.

8. D.J. Lloyd, "Aspects of Fracture in Paniculate Reinforced Metal Matrix Composites," Ada Metallurgica Et Materialia, 39 (1991), 59-71.

9. H.J. Qao, B.H. Ji, I.L. Jager, E. Arzt, and P. Fratzl, "Materials become insensitive to flaws at nanoscale: Lessons from nature," Proceedings of the National Academy of Sciences of the United States of America, 100 (2003), 5597-5600.

10. Z. Zhang and D.L. Chen, "Consideration of Orowan strengthening effect in paniculate-reinforced metal matrix nanocomposites: A model for predicting their yield strength," Scripta Materialia, 54 (2006), 1321-1326.

11. K.S. Tun and M. Gupta, "Improving mechanical properties of magnesium using nano-yttria reinforcement and microwave assisted powder metallurgy method," Composites Science and Technology, 67 (2007), 2657-2664.

12. S. Valdez, B. Campillo, R. Perez, L. Martinez, and A. Garcia, "Synthesis and microstructural characterization of Al-Mg alloy-SiC particle composite," Materials Letters, 62 (2008), 2623-2625.

13. M. Goken and H.W. Hoppel, "Tailoring Nanostructured, Graded, and Particle-Reinforced Al Laminates by Accumulative Roll Bonding," Advanced Materials, 23 (2011), 2663-2668.

14. X.C. Li, Y. Yang, and X.D. Cheng, "Ultrasonic-assisted fabrication of metal matrix nanocomposites," Journal of Materials Science, 39 (2004), 3211-3212.

15. Y. Yang, J. Lan, and X.C. Li, "Study on bulk aluminum matrix nano-composite fabricated by ultrasonic dispersion of nano-sized SiC particles in molten aluminum alloy," Materials Science and Engineering A-Structural Materials Properties Microstructure and Processing, 380 (2004), 378-383.

16. Y.M. Youssef, R.J. Dashwood, and P.D. Lee, "Effect of clustering on particle pushing and solidification behaviour in TiB2 reinforced aluminium PMMCs," Composites Part A-Applied Science and Manufacturing, 36 (2005), 747-763.

17. R.S. Mishra and Z.Y. Ma, "Friction stir welding and processing," Materials Science & Engineering R-Reports, 50 (2005), 1-78.

18. Z.Y. Ma, "Friction stir processing technology: A review," Metallurgical and Materials Transactions A-Physical Metallurgy and Materials Science, 39A (2008), 642-658.

19. R.S. Mishra, Z.Y. Ma, and I. Charit, "Friction stir processing: a novel technique for fabrication of surface composite," Materials Science and Engineering a-Structural Materials Properties Microstructure and Processing, 341 (2003), 307-310.

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20. Y. Morisada, H. Fujii, T. Nagaoka, and M. Fukusumi, "MWCNTs/AZ31 surface composites fabricated by friction stir processing," Materials Science and Engineering A-Structural Materials Properties Microstructure and Processing, 419 (2006), 344-348.

21. H. Choi, M. Jones, H. Konishi, and X.C. Li, "Effect of combined addition of Cu and aluminum oxide nanoparticles on mechanical properties and microstructure of Al-7Si-0.3Mg alloy," Metallurgical and Materials Transactions A-Physical Metallurgy and Materials Science, DOI: 10.1007/sl 1661-011-0905-7 (2011).

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF CORE-SHELLED NANOPARTICLES OF CARBON-COATED NICKEL ON MAGNESIUM

Yi Sun1, Hongseok Choi2, Hiromi Konishi3, Vadim Pikhovich4, Robert Hathaway4, and Xiaochun Li2

'Materials Science Program, University of Wisconsin-Madison, Madison, WI, USA 2Mechanical Engineering Department, University of Wisconsin-Madison, Madison, WI, USA

'Department of Geoscience, University of Wisconsin-Madison, Madison, WI, USA 4Oshkosh Corporation, Oshkosh, WI, USA

Keyword: Magnesium, Metal Matrix Nanocomposite, Core-Shelled Nanoparticles, Mechanical Properties.

Abstract

This paper is to investigate the effect of core-shelled nanoparticles (carbon-coated nickel) on magnesium. With about 4.9wt% of C-coated Ni nanoparticles dispersed by ultrasonic cavitation into pure Mg, the grain size of cast magnesium was refined markedly from more than 450 urn for pure Mg to 25um for Mg nanocomposites. Mechanical property of Mg nanocomposites was also enhanced significantly. The microstructure was examined by Polarized Light Microscopy, SEM, and TEM. TEM study indicates that about 40% of carbon-coated Ni nanoparticles remain intact and are captured inside grains, while the other core-shelled nanoparticles dissolved and formed Mg-Ni intermetallic phase near the grain boundary. Moreover, pure Ni and graphite nanoparticles were also added to pure magnesium to conduct comparison study. Mg with same amount of Ni and C can only produce a much larger average grain size, about 70 um, and lower mechanical properties.

Introduction

Polycrystalline pure Mg exhibits extraordinary high damping properties at engineering stress levels and it will be a prominent material in the control of vibration and noise'1'. The low density of magnesium is also favorable for its application in the automotive and aerospace industry. In the automotive industry, magnesium alloys have been considered one of strong candidate materials for lightweight vehicles with lower fuel consumption and lower carbon dioxide emission. Magnesium alloys will be used for components in all modules of the vehicle (drive train, interior, body and chassis)'2''4'. The unique properties of magnesium alloys, such as light weight, high specific stiffness, and high damping characteristics would be also advantageous for light-weight aerospace structuresI3'.

However, the low strength of pure magnesium limits its widespread applications in industry. Therefore, much research work has been focused on how to enhance mechanical properties of magnesium'1"5'. Grain refinement is often used to enhance service performance of cast products. For aluminum-free magnesium alloys, Zr is very effective as a commercially available grain refiner'8'.

Metal matrix nanocomposites (MMNCs) are attractive in numerous applications because of their improved mechanical properties,6, ' '' n\ Recently, efforts have been taken to develop magnesium matrix nanocomposites due to their higher strength, low density, and higher microhardness.

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For solidification processing of MMNCs, it is crucial to achieve good nanoparticle dispersion and capture. Previous studies'"'2' have proved that high-intensity ultrasonic waves with strong transient cavitation and acoustic streaming could distribute and disperse nanoparticles into magnesium alloy melts successfully, and have produced high-performance nanoparticles reinforced magnesium matrix nanocomposites. However, little research has been conducted to use core-shelled nanoparticles to reinforce Mg alloys during solidification processing. Core-shelled nanoparticles could provide unique combinations of materials properties for nanoparticle interactions with magnesium during solidification processing.

In this study, core-shelled carbon coated nickel nanoparticles will be used to reinforce magnesium to investigate the effect of core-shell nanoparticles on magnesium. Cast Mg/C-coated Ni nanocomposites will be fabricated using a high-intensity ultrasonic wave technique In order to study the effect of Ni addition on mechanical properties and microstructure of pure Mg and for comparison purpose, various content of nickel will also added to Mg melt under same experimental condition To understand carbon coating effect from carbon-coated nickel nanoparticles, graphite nanoparticles will also be added magnesium with nickel addition.

Experimental Procedure

Figure 1 shows the schematics of experimental set-up for the ultrasonic-cavitation based dispersion processing of C-coated Ni nanoparticles (20nm diameter with 2nm graphite coating, as shown in Fig.2) in the molten Mg.

Figure 1. Schematics of experimental set-up for ultrasonic processing.

The system consists of a resistance heating furnace to melt the alloys, a nanoparticle feeding system, gas protection system and an ultrasonic processing system. The mild steel crucible with a diameter of 38.1 mm and a height of 76.2 mm is used for melting, ultrasonic processing and casting. Ultrasonic processing was conducted using a Sonicator 3000 (Misonix Inc., USA) with a niobium ultrasonic probe, which is 12.7 mm in diameter and 92 mm in length.

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Figure 2. TEM image of the C-coated Ni nanoparticles

About 130g of pure Mg was first melted in a mild steel crucible under protective protected by CO2+0.75% S Ff, shielding gas. The tip of the niobium ultrasonic probe was inserted about 12.7 mm in depth into the melt of 700 °C. Ultrasonic vibration with a frequency of 20 kHz and amplitude (peak-to-peak) of 60 um was generated from the transducer, and then 4.9wt% (4.7wt%Ni+0.2wt%C-coating, about 1.0 vol.% total) C-coated Ni nanoparticles were added into the magnesium melt using a small-scale nanoparticle feeding system during ultrasonic processing as shown in Fig. 3. This current system (a Free 3-Axis Small-Scale of Nanoparticle Feeding System) has slides and a lock framing in 3-axis, allowing to accurately position the end tip of feeding tube inside the melt, and also to easily resume the previous position in the next experiment. As shown in Fig.3, the feeding system consists of an auger, a steel feeding tube, a motor, a hopper with a cover lid and inert Ar gas tube, and a frame. The stepper motor was used to drive the auger and dose nanoparticles into the molten magnesium with a controlled feeding rate. The auger and tube allow for a gradual introduction of nanoparticles directly into the cavitation zone under the niobium tip. The hopper has a cover lid with the inert Ar gas tube to obatin a protective environment for nanoparticles. The replaceable steel feeding tube ensures easier cleaning and maintenance.

Figure 3. Small-scale of nanoparticle feeding system

The melt was ultrasonically processed for 20 minutes. After the ultrasonic processing, the ultrasonic probe was lifted out of the melt and the temperature of the melt was increased to 725 °C for pouring. The melt was cast into a steel permanent mold preheated to 350 °C. Two standard flat tensile specimens with a gage dimension of 6.35 mm x 6.35 mm and a gage length

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of 38.1 ram were obtained after each casting. To obtain various content of nickel (such as 2.8 wt.%, 4.7wt% and 4.9wt.%) in Mg for comparison, pure Ni wires (0.5mm diameter, 99.98% purity) were added into the melt manually and used mechanical stir to disperse and form Mg-Ni alloy samples. Graphite nanoparticles were also fed through the feeder into the Mg melt when making samples of Mg-Ni-carbon nanoparticles. These Mg-Ni alloy samples were also cast under same experimental conditions as those for Mg nanocomposites.

Mechanical properties of the samples were determined using a tensile testing machine (SINTECH 10/GL, MTS, USA) with a crosshead speed of 5.08 mm/minute. An extensometer with a 25.4 mm gage length was employed for an accurate measurement of yield strength. When the strain reached 0.2%, the tensile testing machine paused temporarily so that the extensometer can be removed. Final elongation was measured manually after putting the two tested sample pieces back together according to ASTM B557-06. For microstructural characterization, polished samples were lightly etched using an acetic-picric etchant. Optic microscope and polarized light microscope were used for metallographic analysis at this time. Scanning Electron Microscopy (SEM, LEO 1530) and Scanning/Transmission Electron Microscopy (S/TEM, FEI TITAN-80-200) were also utilized for a detailed analysis.

Experimental Results and Discussion

The total Ni content in Mg+4.9wt% C-coated Ni nanocomposite was confirmed by Metallurgical Associates, Inc. using wet chemistry method. The result shows that Ni content is 4.74wt% (6.48g) which means that most of the carbon-coated nickel nanoparticles were added into the matrix successfully.

The optical micrographs of the Mg-Ni alloys and nanocomposite samples are shown in Fig. 4. Both of the Mg+2.8wt%Ni and the Mg+4.9wt%Ni samples consist of the dendrite a-Mg (bright area) and the eutectic phase (dark area), which is fine lamellar shaped a-Mg+Mg2Ni . Suprisingly, there is also eutectic phase in the Mg+4.9wt%C-coated Ni nanocomposite samples. It suggests that the carbon shell on some core-shelled nanoparticles was damaged so that nickel dissolved into molten Mg and forming Mg2Ni during the processing. It is obvious that the content of eutectic phase in the Mg+4.9wt% Ni alloy sample is higher than that in the Mg+4.9wt% C-coated Ni nanocomposite. By using imaging processing techniques, the average area percentages of the eutectic phase in Mg+4.9wt% Ni alloy, Mg+2.8wt% Ni alloy and Mg+4.9wt% C-coated Ni nanocomposite were determined to be 24.9%, 14.2% and 14.0%, respectively. It suggests that about 2.8wt% nickel was dissolved into magnesium during processing for the Mg+4.9wt%C-coated Ni nanocomposites, leaving only about 2.1 wt.% (~0.4vol.%) core-shelled nickel nanoparticles intact in Mg matrix. As shown in Fig.4e, some clusters of C-coated Ni nanoparticles were observed in the optical micrographs. These nanoparticles formed some micro clusters and some were inside the a-Mg grains. In the Mg+4.7wt%Ni+0.2wt%graphite nanoparticle sample, the cluster of graphite particles were also observed but most of them were pushed to the intermetallic phase.

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Figure 4. Optical micrographs of (a) Pure Mg (b) Mg+2.8wt% Ni alloy(c) Mg+4.9wt% Ni alloy (d) Mg+4.7wt%Ni+0.2wt%C (e) Mg+4.9wt% C-coated Ni nanocomposite

Average grain size was measured using the linear intercept method from polarized micrographs of the samples (as shown in Figure 5) according to ASTM E 112-96. Figure 6 shows the average sizes of the samples. For the Mg-Ni alloy samples, the size of primary a-Mg grain decrease (from 462um to 145um) with the increase of Ni content. The graphite nanoparticles in the Mg+4.7wt%Ni+0.2wt%C sample introduce further grain refinement to obtain a smaller grain size, 70um. However, the Mg+4.9wt% C-coated Ni nanocomposites provide the most significant grain refinement among all samples with a grain size of 25um. Additional grain refinement in the case of core-shelled nickel reinforced nanocomposite samples can be primarily attributed to the coupled effect of: (i) capability of fine second phase nanoparticles to nucleate magnesium grains, and (ii) restricted growth of magnesium grains as a result of pinning by the presence of finer reinforcement C-coated Ni nanoparticles. This dispersed nanoparticles also help to refine/modify the eutectic Mg-Mg2Ni phase in the Mg+4.9wt% C-coated Ni nanocomposite sample, leading to good ductility.

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Figure 5. Polarized light micrographs of (a) Pure Mg (b) Mg+2.8wt% Ni alloy(c) Mg+4.9wt% Ni alloy(d) Mg+4.9wt% C-coated Ni nanocomposite (e) Mg+4.7wt% Ni wire+0.2wt% C

nanoparticles

Figure 6. Average grain sizes of Mg-Ni alloys and nanocomposites

The SEM images (Fig.7) of the Mg+4.9wt%C-coated Ni nanocomposite prove that some nanoparticles survived in the molten Mg. The EDS spectrum of some particles in the Mg matrix (Fig.7c) showed Ni peak which confirm that the C-coated Ni nanoparticles were captured in the

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grain. This capture process would also restrict the growth of magnesium grains. However, in the Mg+4.7wt%+0.2wt%C sample, almost of all the graphite nanoparticles were pushed to the eutectic phase and no capture was observed. This further suggests that carbon-coated Ni nanoparticles provide a powerful means for significant grain refinement.

Figure 7. SEM images of (a) Mg+4.9wt% Ni alloy and (b) Mg+4.9wt%C-coated Ni (c) EDS analysis of Mg+4.9wt%C-coated Ni nanocomposite

The TEM images (Fig. 8a) clearly show that some carbon-coated nickel nanoparticles were pushed near the eutectic phase while still captured by the a-Mg (Fig.8b). Only the nanoparticles with thick C-coating (more than 3nm) survived the processing while these core-shelled nanoparticles with thin coating and/or weaker bonding dissolved into Mg to form the eutectic phase.

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Figure 8. TEM images of (a) C-coated Ni nanoparticles in the eutectic phases and (b) C-coated Ni nanoparticles captured by Mg matrix

Fig. 9 shows the tensile testing result of all samples. For the Mg-Ni alloy samples (Mg+2.8wt%Ni wire, Mg+4.9wt%Ni wire), the yield strength and tensile strength increase with the increase of Ni content while the ductility decreases due to increased content of brittle Mg2Ni phase. This result agrees well with what was reported by Wan et al'10'. As for the Mg+4.9wt% C-coated Ni nanocomposite, the yield strength is enhanced by 193% and 42% when compared to the pure Mg sample and the Mg+2.8wt% Ni sample. The yield strength of Mg+4.9wt% C-coated Ni nanocomposite is also 21% higher than that of Mg+4.7wt% Ni wire+0.2wt% graphite nanoparticle sample. Due to the addition of Ni, the Mg-Ni alloy samples all have lower ductility than the pure Mg (6.08%). However the Mg+4.9wt% C-coated Ni nanocomposites maintain relatively high ductility (5.84%). The elongation of the Mg+4.9wt% C-coated Ni nanocomposite slightly decreases from 6.08% to 5.84% possibly due to some microclusters of nanoparticles. It suggests that C-coated Ni nanoparticles effectively enhance yield strength, tensile strength while maintaining good ductility for Mg.

Figure 9. Mechanical properties of Mg-Ni alloy and Mg-C-coated Ni nanocomposites.

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In order to undertsand stengthening mechanisms from C-coated Ni nanoparticles, the Mg+4.9wt% C-coated Ni nanocomposite was compared with the Mg+2.8wt% Ni alloy, which has almost the same content of Mg-Mg2Ni eutectic pahse. Grain refinemnt effect for strenghtening was first obtained by using the Hall-Petch equation,

^Hall-Petch = ky ' (dm2 - dc

2 ) , (1)

Where ky is the Hall-Petch constant (133MPa-um"2)[6', d™ is the average grain size of the Mg+2.8wt% Ni alloy sample, dc is the average grain size of the Mg+4.9wt% C-coated Ni nanocomposite. The predicted yield strength enhancement from Hall-Petch strengthening is 15.55MPa, which is less than the total enhancement in the experiment result (25 MPa between pure Mg and Mg-lvol% C-coated Ni nanocomposite). Orowan strengthening must also be considered by

t±°nrnwan — r î T *" 7T% \*-)

\2Vp)

2b'

Where Vp is the volume fraction of the remaining nanoparticles (0.4 vol%, derived from the area fraction of the eutectic phase), Gm is the shear modulus of the matrix ( 11.60 GPa), dp is the diameter of the nanoparticles (20 run), b is the Burgers vector of the matrix (0.35 nm). The calculation predicted that the yield strength enhancement due to Orowan strengthening is 22.21 MPa. Thus, the total yield strength enhancement predicted by the combination of Hall-Petch and Orowan strengthening is to be 37.76 MPa, which is higher than the experimental data (25 MPa). This is possibly due to some microclusters of the C-coated Ni nanoparticles and some nanoparticles were pushed to grain boundaries.

Conclusions

4.9wt.% C-coated Ni nanoparticles with a diameter of 20 nm were successfully fed into magnesium melt using the upgraded free 3-axis small scale of nanoparticle-feeding system. The yield strength of Mg+4.9wt% C-coated Ni nanocomposite reaches 85MPa and the nanocomposite maintains a good ductility as well. Polarized light micrographs showed a significant grain refinement by the addition of C-coated Ni nanoparticles. The average grain sizes of pure Mg, Mg+4.9wt% Ni alloy, Mg+4.7wt% Ni+0.2wt% C and Mg+4.9wt% C-coated Ni nanocomposites were 440 urn, 140 urn, 70um and 25 um, respectively. Experimental results suggest that both of the C-coated Ni nanoparticles and the graphite nanoparticles worked as heterogeneous nucleation agents. However, the C-coated Ni nanoparticles introduce more significant grain refinement possibly due to the restricted growth of grains by nanoparticle captures. Optical, SEM and TEM images indicate that only about 40% of the original 4.9wt% C-coated Ni nanoparticles survived and captured by the Mg matrix. These survived core-shelled nanoparticle normally had a thicker carbon coating (3~4nm). The eutectic phase Mg-Mg2Ni is refined in the Mg+4.9wt% C coated Ni nanocomposite due to grain refinement and nanoparticles restricted growth of eutectic phases. Theoretical study suggests that the yield strength enhancement is due to both Hall-Petch strengthening and Orowan strengthening. The difference between the enhancement prediction and experimental result is likely caused by some micro-clsuters of the C-coated Ni nanoparticles.

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Acknowledgements

This work is sponsored by National Institute of Standard and Technology through its Technology Innovation Program.

Reference

[I] X.S. Hu et al., "A study of damping capacities in pure Mg and Mg-Ni alloys," Scripta Materialia, Vol. 52, pp. 1141-1145,2005

[2] H. E. Friedrich, B. L. Mordike, "Magnesium Technology Metallurgy, design Data, Applications," Springer-Verlag Berlin Heidelberg, 2006

[3] H. Furuya, N. Kogiso, S. Matunaga, K. Senda, "Application of magnesium alloys for aerospace structure systems," Materials Science Forum, Vol. 350-351, pp. 341-348, 2000

[4] H. Friedrich, S. Schumann, Research for a "new age of magnesium" in the automotive industry, Journal of Materials Processing Technology, Vol. 117, pp. 276-281, 2001

[5] R. Schaller, Metal matrix composites, "a smart choice for high damping materials, Journal of Alloys and Compounds," Vol. 355, pp. 131-135,2003

[6] M. DeCicco, X.C. Li et al, "Strong, Ductile Magnesium-Zinc Nanocomposites," Metallurgical and Materials Transactions A, Vol. 40A, pp. 3038-3045,2009

[7] Z. Zhang, D.L. Chen, "Consideration of Orowan strengthening effect in particulate-reinforced metal matrix nanocomposites: A model for predicting their yield strength," Scripta Materialia, Vol. 54, pp. 1321-1326, 2006

[8] D. H. Süohn, M. Qian, M. A. Easton etc., "Grain Refinement of Magnesium Alloys," Metallurgical and Materials Transaction A, Vol. 36A, pp. 1679, 2005

[9] X.S. Hu, Y.K. Zhang, MY. Zheng, K. Wu, "A study of damping capacities in pure Mg and Mg-Ni alloys," Scripta Materialia, Vol. 52, pp. 1141-1145, 2005

[10] D.Q. Wan, J.C. Wang et al, "Effect of eutectic phase on damping and mechanical properties of as-cast Mg-Ni hypoeutectic alloys," Transactions of Nonferrous Metals Society of China, Vol. 19,pp.45-49,2009

[II] G. Cao, H. Choi, J. Oportus, H. Konishi, and X. Li, "Study on tensile properties and microstructure of cast AZ91D/AIN nanocomposites," Mater. Sei. Eng. A, Vol 494, p 127-131,2008

[12] H. Choi, N. Alba-Baena, S. Nimityongskul, M. Jones, T. Wood, M. Sahoo, R. Lakes, S. Kou, and X. Li, "Characterization of hot extruded Mg/SiC nanocomposites fabricated by casting," J. Mater. Sei., Vol 46, p 2991-2997, 2011

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

IN SITU COMPOSITE OF (Mg2Si)/Al FABRICATED BY SQUEEZE CASTING

Hiiseyin Murât Lus1, Gökhan Ozer', Kerem Altug Guler1

1 Yildiz Technical University, Department of Metallurgical and Materials Engineering, Davutpasa, Istanbul, 34210, TURKEY

Keywords: Squeeze casting, Mg2Si/Al composites, in situ

Abstract

An in situ formed Mg2Si/ Al-Si-Cu reinforced metal matrix composite fabricated by squeeze casting are investigated. It is showed that primary Mg2Si crystals are formed by adding pure Mg into hypoeutectic Al-Si-Cu alloy (A3 80). In order to increase castability and to obtain better properties squeeze casting technique is used. The results show that, the average size of primary Mg2Si particulates decreases from 87um to21 urn. Furthermore, the average porosity values of cast samples are significantly decreased from %8.7 to %0.5 with the application of 30 MPa pressure during solidification.

Introduction Aluminium based particulate metal matrix composites (PMMCs) offer attractive performance in aerospace and automotive applications because of their improved properties such as excellent wear resistance, low density and higher mechanical strength [1], Although, there are several fabrication techniques, today metal matrix composites (MMCs) are generally produced either by casting or powder metallurgy methods [2]. In the past years, it was reported [3] that the worldwide MMC market in 1999 accounted for 2500 metric tons valued at over $100 Million and is expected to exceed $282 Million by 2015 [4]. However, poor wetting of reinforcement phase by molten metal, non-uniform distribution of ceramic particulates and relatively high-cost fabrication techniques prevents the large scale industrial production of ex-situ PMMCs in automotive and high-safety applications [5]. Over the past twenty years in situ Mg2Si/Al based Metal Matrix Composites have emerged as an important class among various Al-based composites, since (a) low density and high melting point of Mg2Si particles and (b) elimination of wetting problems [6]. Because of these advantages, many researchers have focused on investigating Al based- Mg2Si composites [7, 8], However the properties of these composites are limited by the low ductility due to the existence of hard and fragile primary Mg2Si and Si [9]. In recent years, many works were focused on the improvement of microstructure and mechanical properties by adding alloying elements and modifiers [10, 11]. On the other hand, it is well known fact that cast aluminium alloys are susceptible to have relatively higher porosity values which may deter its use in high-safety applications. Besides, Mg2Si/Al-Si-Cu alloys have relatively longer freezing range caused porosity content in Al-Mg2Si composites. Therefore, it is necessary to develop feasible alternative fabrication techniques to improve the quality of cast components. Squeeze casting also known as liquid-metal forging, squeeze forming, extrusion casting is a casting process in which liquid metal is solidified under direct pressure in permanent die cavity [12]. Squeeze casting has several advantages over convantional casting methods such as obtaining very fine microstructure with near-net shape and elimination of gas and shrinkage

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porosities, etc. This work presents the effect of applied pressure on microstructure and porosity content of Mg2Si/Al-Si-Cu reinforced composite fabricated by direct squeeze casting process.

Experimental procedure Sample preparation

Commercially purchased Al-8.7%Si-3%Cu alloy and pure Mg (ingot) were used as starting materials. About 850 g. Al-Si-Cu alloys were melted in SiC crucible at electrically heated furnace. About 150 g magnesium preheated at 300 °C were added into the melt at 680-700 °C. After holding for 15 min at this temperature, the in situ Mg2Si/ Al-Si-Cu composite melts were poured into a steel die to produce ingots of 023x100 mm.

Table 1 Chemical Composition of in situ Mg2Si/Al-Si-Cu composites (wt. %)

Element

Content (wt %)

Si

7,83

Fe

0.284

Cu

2.78

Mg

15,62

Mn

0.308

Al

Rest.

In order to compare microstructural change, same alloy was squeeze cast with 30 MPa pressure during solidification. The composite test specimens were fabricated by squeeze casting using cylindrical H13 tool steel dies. The lower die was mounted on a supporting unit and the upper die was attached to a hydraulic press ram (Fig 1). The steel dies were coated with Zyp Coatings Boron Nitride Aerosol and then air dried. Prior to squeeze casting, the dies were preheated to 150 °C with an electric resistance furnace. Then the melt poured into steel lower mold and squeezed with 30 MPa. The final alloy composition is listed in Table 1. Metallographic specimens were prepared by standard polishing and etching procedure with Kellers reagent. Microstructure and phase of the in situ composite were investigated by using optical microscope (Leica Qwin-Image analyzer). The amount of porosity was measured using approximately 10x10x10 mm3 specimens for each casting method with the Archimedes' principle as described by Taylor et al [13]. To eliminate the other effects of parameters, four specimens were taken from the different region from cast pieces. Hardness was determined by using standart Brinell hardness equipment with 62.5 kg load from the average of five readings from each sample.

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Figure 1 a) Die-set used in the squeeze casting process and b) cross section of lower die

Results and discussion Due to wide freezing range of this alloy, the primary Mg2Si particles will become rather coarse in the gravity die casting alloy. On the other hand, with the increasing of cooling rate, decreasing of the average size of Mg2Si phase is expected. Fig.2 shows the optical microscope microstructures of of (a) gravity die cast and (b) 30 MPa squeeze cast composites. The results reveal that both gravity and squeeze in situ composites consist of dark polygonal shaped Mg2Si, grey a-Al and very little amount CUAI2. In detailed, in both gravity and squeeze cast composites, primary Mg2Si is identified as dark faceted polygon shaped morphology with a hole filled with a-Al with with different size Fig 3. It is observed that with the application of 30 MPa pressure, the size of in situ Mg2Si decreased from 87 um to 21 urn. Zhang et al reported that, Mg and Si have very limited diffusion in liquid when the cooling rate is high during solidification [14]. According to the Clausius-Clapeyron equation, freezing point of metals and alloys increase with the application of pressure resulting in larger degrees of undercooling in the melt thus increases nucleation frequency [15]. Elimination of air gap between the alloy and the die wall also increaes heat-transfer coefficients and cooling rate considerably. Consequently, formation Mg2Si primary particles with a smaller size can be explained by the increase of undercooling degree resulting from increase in melting point and enhanced heat-transfer coefficient.

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Figure 2 Microstructures of Mg2Si/Al-Si-Cu in situ composites with (a) 0 Mpa, (b) 30 MPa

Figure 3 SEM microstructure of squeeze cast Mg2Si/Al-Si-Cu composite

It is well known that Al composites freeze over a wide range of temperature are prone to have shrinkage or gas porosity. One of the main reasons for the existence of shrinkage porosities in cast parts is ineffective feeding systems in cast processes. High pressure during solidification can effectively improve the feeding capacity of castings which hinders the shrinkage porosity formation [16]. With the application of pressure during solidification, gas solubility in the melt increases and the nucleation of gas bubbles becomes more difficult. In Fig. 2a porosity formation

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can be seen easily in the alloy when the alloy is cast with no applied pressure (bottom of the picture). On the other hand, gas and shrinkage porosities and/or a combination of them formed in the gravity die cast were completely disappeared in squeeze cast sample. Effect of applied pressure on the porosity percent is shown in Fig. 4.

Figure 4 Effect of pressure on the porosity percent of Mg2Si/Al-Si-Cu in situ composites

Conclusion

The effects of pressure during solidification on microstructure and volume fraction of gas and/or shrinkage porosities in in situ composite of Mg2Si/Al-Si-Cu are investigated. The following conclusions may be drawn from this study that:

1. In situ composites of Mg2Si/Al-Si-Cu reinforced composite alloy can successfully be cast using direct squeeze casting process. Primary Mg2Si particle was refined and their sizes decreased from -87 um and to -21 urn with the application of 30 MPa pressure during solidification.

2. The applied pressure in squeeze casting significantly decreases the volume fraction of porosity in cast samples from %8.7 to %0.5.

Acknowledgement

The authors wish to thank BCACT (Balkan Centre for Advanced Casting Technologies) for their help in experiments.

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References

[1], T.P.D. Rajan, K. Narayan Prabhu , R.M. Pillai, B.C. Pai, "Solidification and casting/mould interfacial heat transfer characteristics of aluminum matrix composites" Composites Science and Technology, Vol. 67 (1) ( 2007), 70-78.

[2] J.W. Kaczmar, K. Pietrzak, W. Wlosihski, "The production and application of metal matrix composite" Journal of Material Processing Technology, Vol. 106 (1-3) (2000), 58-67.

[3] D. B. Miracle, "Metal Matrix Composites - From Science to Technological Significance", Compos. Sei. Technol, 65 (2005), 2526-2540.

[4] Metal Matrix Composites: A Global Strategic Business Report, Global Industry Analysts, Inc. (2011).

[5] J. Zhang, Z. Fan, Y. Q. Wang, B. L. Zhou, "Microstructural evolution of the in situ Al-15wt.%Mg2Si composite with extra Si contents", Scripta mater. , 42 (2000), 1101-1106.

[6] G. Frommeyer, S. Beer, K. Von Oldenburg, "Microstructure and properties of as-cast

intermetallic Mg2Si-Al alloys" Zeitschrift fir Metallkunde, 81 (1990), 809-815.

[7] G. Frommeyer, S Beer and K. Von Oldenburg, "Microstructure and mechanical properties of mechanically alloyed intermetallic Mg2Si-Al alloys" Zeitschrift für metallkunde, 85 (1994), 372-377.

[8] B.S.S. Daniel, V.S.R. Murthy, G.S. Murty, "Metal-ceramic composites via in-situ methods", Journal of Materials Processing Technology, 68 (2) (1997), 132-155.

[9] Q.D. Qin, Y.G. Zhao , K. Xiu, W. Zhou, Y.H. Liang, "Microstructure evolution of in situ Mg2Si/Al-Si-Cu composite in semisolid remelting processing", J. Mater Sei Eng A, 407 (2005), 196-200.

[10] Y.G. Zhao, Q.D. Qin, W. Zhou, Y.H. Liang, "Microstructure of the Ce-modified in situ Mg2Si/Al-Si-Cu composite", Journal of Alloys and Compounds, 389 (1-2) (2005), L1-L4.

[11] Q.D. Qin, W.X. Li , K.W. Zhao, S.L. Qiu, Y.G. Zhao, "Effect of phosphorus on microstructure and growth manner of primary Mg2Si crystal in Mg2Si/Al composite", J. Mater Sei Eng A, 527(2010), 2253-2257.

[12] L. J. Yang, "The effect of casting temperature on the properties of squeeze cast aluminium and zinc alloys", Journal of Materials Processing Technology, 140 (1-3) (2003), 391-396.

[13] R.P. Taylor, S.T. McClain and J.T. Berry, "Uncertainty Analysis of Metal- Casting Porosity Measurements Using Archimedes' Principle", International Journal of Cast Metals, 11 (1999), 247-257.

[14] J. Zhang, Z. Fan, Y. Q. Wang, B. L. Zhou, "Microstructural evolution of the in situ Al-15wt.%Mg2Si composite with extra Si contents", Scripta mater. , 42 (2000), 1101-1106.

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[15] M.R. Ghomashchi, A. Vikhrov, "Squeeze casting: an overview", Journal of Materials Processing Technology. 101 (2000), 1-9.

[16] A. Maleki, B. Niroumand, A. Shafyei, "Effects of squeeze casting parameters on density, macrostructure and hardness of LM13 alloy", Materials Science and Engineering A 428 (2006), 135-140.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

OPTIMIZATION OF TENSILE STRENGTH OF FRICTION STIR WELDED Al-(10 tol4 wt.%) TiB2 METAL MATRIX COMPOSITES

Santhiyagu Joseph Vijay1, Nadarajan Murugan2, Siva Parameswaran3

1 Assistant Professor, School of Mechanical Sciences, Karuny a University, Coimbatore, Tamil Nadu, India-641114. viiavioseph@karunva.edu

2 Professor, Department of Mechanical Engineering, Coimbatore Institute of Technology, Coimbatore, Tamil Nadu, India - 614 014. Murugan@cit.edu.in

2 Professor, Department of Mechanical Engineering, Texas Tech University, Lubbock, Texas, TX 79409-1021, USA. siva.parameswaran@ttu.edu

Keywords: Metal Matrix Composites, Friction Stir Welding, Tensile Strength, Process Parameters, Optimization.

Abstract Metal Matrix Composites (MMCs) play a vital role in replacing many structural materials

due to their superior mechanical and metallurgical properties. Processing those composites has always been a major factor influencing their application. AI-TiB2 MMCs are used as structural members in the marine industry. Friction Stir Welding (FSW) has revolutionized the process of joining those composites.

In this paper, an attempt has been made to friction stir weld the AI-TiB2 MMCs and to develop a regression model for predicting the tensile strength of the weldment. The process parameters considered for FSW are tool rotation speed, tool traverse speed, axial load and weight percentage of TiB2 in Al matrix. The regression model is used to optimize the process parameter using Desirability Optimization Methodology to improve the tensile strength of the FS welded Al-TiB2 composites. The effects of process parameters on the tensile strength of the welded composites are analyzed and presented.

Introduction Particle reinforced metal matrix composites (MMCs) have already received greater attention

from aerospace, automobile and marine industries as a substitute for their monolithic counterparts due to their higher stiffness, superior strength, improved resistance to wear and low coefficient of thermal expansion [1]. The applications of these MMCs are limited due to non-existence of standardized procedure to process them like machining or welding. The presence of reinforced ceramic particles in the MMCs especially reduces their weldabilty. Attempts were made to weld them using fusion welding processes which resulted in porosity, coarse microstructure, segregation and decomposition of ceramic particles and formation of brittle intermetallic compounds. The high temperature during fusion welding increases the tendency of ceramic particles to react with aluminum matrix. Achieving homogeneous distribution of ceramic particles in the weld zone was seldom possible. The mechanical properties of the joints were also poor [2-5]. To overcome these problems of fusion welding process, researchers have proven that these materials can be welded using friction stir welding (FSW) process.

The invention of FSW process by TW1 in 1991 has been a revolution in the field of materials joining [6 -13]. The mechanical properties of the joints made by FSW of MMC depends on various parameter such as tool rotation speed, tool traverse speed, axial force and weight % of reinforcement particles in the aluminum matrix. Many researchers have already established empirical relations using the central composite rotatable design to conduct experiments and correlate the effect of various FSW process parameters to the mechanical properties of the weldment [14-16].

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In this paper, an attempt has been made to find the influence of FSW process parameters on the tensile strength of FS welded AA6061-TiB2 MMCs. The MMCs were produced using in-situ stir casting process containing weight percentage of TiB2 varying from 10% to 14%. Though there are various methods to produce MMCs, the salt route was chosen due to the advantages like easy control of phases, less contamination and possibility of bulk casting [17-21]. Experiments were conducted on the Al-TiB2 cast MMC plates according to central composite rotatable design. The empirical relationship between FSW process parameters and tensile strength was developed and optimized to achieve better tensile strength in the weldment.

Experimental Work FSWofAl-TiB2MMCs

Al-TiB2 composites were fabricated using in-situ stir casting process. FSW of the produced Al-TiB2 MMCs were carried out in the indigenously developed FSW machine (Fig. 1). The welding specimen of size 100 mm x 50 mm x 6 mm obtained from the stir casting using WEDM machine. The FSW process parameter chosen are tool rotation speed, tool traverse speed, axial load and % reinforcement of TiB2 in Aluminum matrix. A non-consumable square pin profiled tool made of High Carbon High Chromium Steel was used for welding. Welding trials were carried out on the specimen to fix the working range of the above said process parameters. The limits of these parameters were derived from the trial welds so that the joints are free from any visible external defects. The important factors influencing the tensile properties of FSW joints and the working ranges of those factors for Al-TiB2 MMC are presented in Table I.

Figure 1. Friction Stir Welding Machine Figure 2. Typical Tensile Specimen Table I. Coded Values and Working Range of FSW Process Parameters.

Process Parameters

Tool rotation speed Tool traverse speed

Axial load Reinforcement

Notation

N S F R

Unit

rpm mm/min

kN %

(-2) 1300 60 9 10

(-1) 1475 75

11.75 11

Levels (0)

1650 90

14.75 12

(1) 1825 105

17.25 13

(2) 2000 120 20 14

A four factors, five levels central composite design matrix was derived to conduct the experiments. Table II shows the 31 sets of coded conditions used to form the design matrix. As prescribed by the design matrix, thirty-one joints were fabricated. According to ASTM E08 Standards, three tensile specimens were fabricated from each weldment using a power hacksaw and then machined to the required size, as shown in Fig. 2. Tensile strengths of the FSW joints were evaluated by conducting testing in an UTM, and the average of three results is presented in Table II.

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Deri

Run No.

1 2 3 4 5 6 7 8 9 10 11 12 13 14 15

Table I Design Matrix

N -1

-1

-1

-1

-1

-1

-1

-1

S -1 -1 1 1

-1 -1 1 1

-1 -1 1 1

-1 -1 1

F -1 -1 -1 -1

-1 -1 -1 -1

R -1 -1 -1 -1 -1 -1 -1 -1

vine a Mathematical Model

. Design

UTS (MPa)

285 287 290 283 289 290 291 280 300 302 311 304 308 310 314

Matrix and Experimental Results.

16 17 18 19 20 21 22

23 24 25 26 27 28 29 30 31

1 -2 2 0 0 0 0 0 0 0 0 0 0 0 0 0

1 0 0 -2 2 0 0 0 0 0 0 0 0 0 0 0

1 0 0 0 0 -2 2

0 0 0 0 0 0 0 0 0

1 0 0 0 0 0 0 -2 2 0 0 0 0 0 0 0

310 293 297 292 294 290 291 284 316 296 297 300 290 296 299 302

Based on the process parameters under consideration and the response that is taken into account a mathematical model can be derived for the response as a function of the process parameters i.e., the tensile strength (UTS) of the weldment is a function of tool rotation speed (N), tool traverse speed (S), axial force (F), % reinforcement of TiB2 (R) and is denoted with the following expression. UTS = f (N, S, F, R) (3)

The regression equation with second order polynomial is used to represent the response surface Y and is given by [21, 22] Y = bo + Ib i x i + Ib i ix i

2 + IbijX,xJ (4) The four process parameter can be expressed using the same polynomial equation as given

below. UTS = bo + b, (N) + b2 (S) + b3 (F) + b, (R) + b„ (N2) + b22 (S

2) + b33 (F2) + b44 (R

2) + b,2 (NS) + b,3

(NF) + b14 (NR) + b23 (SF) + b24 (SR) + b34 (FR) (5) Where, bo is the average of responses and bi, b2, ..., b34 are the coefficients of the respective

main and interaction effects of the parameters. All the coefficients were evaluated and tested for their significance at a 95% confidence level by applying a student's t-test using SYSTAT and Design Expert statistical software packages. After determining the significant coefficients, the final model was developed using only those coefficients and is given below. UTS = {297.14 - 0.58N + 0.67S + 1.33F + 9.50R - 0.12N2 - 0.62S2 - 1.24F2 + 1.13R2 - 2.25NS -

0.12NF} MPa (6) The coefficient of determination is denoted as 'r2' and is used to find how close the predicted

and experimental values lie [21]. The r2 value indicates the goodness of fit for the model. In this case, the value of r2 = 0.92, which indicates that only less than 8 % of the total variations are not explained by the model. This indicates, an excellent suitability of the regression model.

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Optimization of FSW Parameters The desirability optimization technique available in the Design Expert software package is

used to optimize the process parameters that are under consideration. The objective function for optimization is the mathematical model which was derived earlier and it is a maximization of Tensile Strength. The constraints are the lower limits and upper limits of the process parameters. The expressions for the mathematical model and constraints are given as below. Maximize (UTS) = 297.14 - 0.58N + 0.67S + 1.33F + 9.50R - 0.12N2 - 0.62S2 - 1.24F2 + 1.13R2 -

2.25NS-0.12NF (7) Bounded to: -2 < (N) < 2; -2 < (S) < 2; -2 < (F) < 2; -2 < (R) < 2. (8) The optimized values for the process parameters are: i) Tool Rotation Speed = 1433 rpm; Tool Traverse Speed = 117 mm/min; Axial Force = 14.4 kN and % Reinforcement of TiB2= 14 %. The optimized UTS = 328.8 MPa.

Results and Discussion Effect of tool rotational speed on Tensile strength

Fig. 3 indicates that the joints fabricated using the higher rotational speed of 2000 rpm exhibit inferior tensile properties. An increase in the tool rotation speed has caused the increase in the heat input and due to which the average grain size is increased. This leads to the lower bonding strength and thus inferior tensile strength. Table III shows the macrograph of various welded specimen fabricated at 1300 rpm, 1650 rpm and 2000 rpm. In the weld nugget, coarsening of precipitates occurs at high speeds when compared to the parent material. Moreover, at high rotational speeds, TiBî particles would suffer more fragmentations. Thus these microstructural variations affected the mechanical properties (Fig. 4) [22]. The smaller grain size and uniform distribution of TiB2 in the weld nugget are the reasons for the higher hardness and superior tensile properties of the joints made at low tool rotation speed of 1300 rpm. Effect of tool traverse speed on Tensile strength

The joints fabricated using lower and higher welding speeds (60 mm/min and 120 mm/min) exhibit significant variations in tensile properties compared to the joints fabricated using a welding speed of 90 mm/min. The joints fabricated at a welding speed of 90 mm/min have higher tensile strength than that of the other joints (Fig. 5). Table 111 shows the macrographs of weld nugget at various welding speeds such as 60, 90, 120 mm/min. At lower welding speeds, grain coarsening occurs due to higher heat input. Welds produced at the lower welding speed contains grains having many sub-boundaries. The grain size increases with increasing heat input during FSW [22], in the weld nugget. Abrupt change in the grain size in the weld nugget is observed in the joints (Fig. 6) fabricated with higher welding speeds. At higher welding speeds, the weld nugget is exposed to higher temperature for a short span of time. Thus at welding speed of 90 mm/min, finer grains are abserved in the weld nugget resulting in superior tensile strength in the joints. Effect of axial force on tensile strength

Fig. 7 shows that the joints fabricated using lower and higher axial forces of 9 kN and 20 kN are lower when compared to the joints fabricated using an axial force of 14.75 kN. At the lower axial force, the elongated grains of the base material are partially changed into equiaxed grains in the weld nugget due to insufficient forging pressure and friction heat. At the higher axial force, grain growth and grain coarsening due to higher heat input and excessive forging pressure may be the reasons (Fig. 8) [22]. The reason for the superior tensile strength in the joints made with the axial force of 14.75 kN may be due to smaller grains formed in the weld nugget at this axial force. Table III shows the macrograph of joints made at various axial loads.

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Figure 3. Effect of Tool Rotation Speed on Tensile Figure 4. SEM image of the weld nugget Strength obtained at 2000 rpm

Effect of TiB? Particle on Tensile Strength Fig.9 shows that the tensile strength of the welded composite increases with increase in the

weight percentage of TiB2 particles. The homogeneous distribution of the TiB2 particles in the weld nugget is one of the reasons for the improved tensile strength when the wt. % of TiB2 increases. It may also be due to the identical distribution of TiB2 particles in the weld nugget and parent material. The microstructure of the weld nugget of joints made with 14% of TiB2 particles (Fig. 10) shows the homogeneous distribution of TiB2 particles.

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r- c T-«- » CL I T j T -i d .t. Figure 6. SEM image of the weld Figure 5. Effect of tool Traverse speed on Tensile Strength B x L i . , ^ ,„„ , . nugget obtained at 120 mm/min

Interaction Effect of Tool Rotation Speed and Tool Traverse Speed on Tensile Strength. Fig. 11 shows the interaction effect of tool rotation speed and tool traverse speed on tensile

strength. From the graph it is evident that at welding speeds below 90 mm/min, the tensile strength increases with increase in tool rotation speed. Whereas for welding speeds above 90mm/min, the tensile strength decreases with respect to the increase in tool rotation speed. At welding speed of 90mm/min, only a negligible variation in the tensile strength is observed. Thus there seems to be pronounced interaction effect between tool rotation speed and tool traverse speed.

„. _ r „ t . . . , _ _ .. Ct ., Figure 8. SEM image of the weld nugget Figure 7. Effect of Axial Force on Tensile Strength ° th ?0kN If

Figure 9. Effect of TiB2 Particles on Tensile Figure 10. SEM image of the weld nugget with Strength 14% reinforcement of TiB2 Particles

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Figure 11. Interaction effect of tool rotation speed Figure 12. Interaction effect of tool rotation and tool traverse speed on UTS speed and axial force on UTS

Interaction Effect of Tool Rotation Speed and Axial Force on Tensile Strength. From Fig. 12 it is evident that when the axial load increases, the tensile strength tends to

increase in the initial stages but decreases as the load is increased to the maximum level. The same pattern is observed for various tool rotation speeds. At 1300 rpm, the curve traces the highest tensile strength values whereas at 2000 rpm, the curve traces the lowest tensile strength values.

Conclusion • A mathematical model has been developed to predict the tensile strength of friction stir

welded Al-TiB2 MMC joints. • Optimization of the FSW parameters was carried to attain maximum tensile strength in the

welded joints. • The optimized values closely match the experimentally determined values. • Response graphs have been drawn to study the direct and interaction effect of welding

parameters on the tensile strength of friction stir welded joints of Al-TiB2 MMC. Acknowledgement

The authors acknowledge the financial support rendered by Naval Research Board, DRDO, Govt. of INDIA, vide funded project; Ref. no. DNRD/05/4003/NRB/85 dt 30.10.2006. The authors are also grateful to the Dept. of Mechanical Engineering, Coimbatore Institute of Technology, Coimbatore, INDIA for extending the facilities of Welding Research Cell to carry out this investigation. The authors also acknowledge Karunya University, INDIA and Dr. V. Balasubramaniam, Director - CEMAJOR, Annamalai University, INDIA for extending the testing facilities. The authors wish to thank Mr, I. Dinaharan, Mr. Palanivel, Mr. S. Gopalakrishnan, Mr. Harrison for the assistance offered in executing the above work.

Reference 1. Jung-Moo Lee et al., "A New Technology for the Production of Aluminum Matrix

Composites by the Plasma Synthesis Method," Metals and Materials, 6 (2000), 389-394 2. Rosso M, "Ceramic and metal matrix composites: Routes and properties," J Mater Process

Technol., 175 (2006), 364-75. 3. Huang RY, Chen SC, Huang JC, "Electron and Laser Beam Welding of High Strain Rate

Superplastic Al-6061/SiC Composites," Metall. Mater. Trans. A, 32 (2001), 2575-84. 4. Lean PP, Gil L, Urefia A, "Dissimilar welds between unreinforced AA6082 and

AA6092/SiC/25p composite by pulsed-MIG arc welding using unreinforced filler alloys Al-5Mg and Al-5Si," J Mater Process Technol., 143 (2003), 846-850.

5. Storjohann D et al., "Fusion and Friction Stir Welding of Aluminum-Metal-Matrix Composites," Metall Mater Trans A, 36 (2005), 3237-47.

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6. Nandan R, DebRoy T, Bhadeshia HKDH, "Recent advances in friction stir welding-process, weldment structure and properties," Prog Mater Sei., 53 (2008), 980-1023.

7. Thomas WM et al., "Friction stir butt welding," International Patent Application No. PCT/GB92/02203, (1991).

8. Dawes CJ, Thomas WM, "Friction stir process welds aluminum alloys," Weld J., 75 (3) (1996), 41-5.

9. Cho JH, Boyce DE, Dawson PR, "Modeling strain hardening and texture evolution in friction stir welding of stainless steel," Mater Sei Eng. A, 398 (2005), 146-63.

10. Chionopoulos SK et al., "Effect of tool pin and welding parameters on friction stir welded (FSW) marine aluminum alloys," Proceedings of the 3rd international conference on manufacturing engineering (ICMEN), Greece, (2008), 307-16.

11. Yingchun Chen, Huijie Liu, Jicai Feng, "Friction stir welding characteristics of different heat-treated-state 2219 aluminum alloy plates," Mater Sei Eng. A, 420 (2006), 21-5.

12. Ceschini L et al, "Effect of friction stir welding on microstructure, tensile and fatigue properties of the AA7005/10 vol.% A1203p composite," Compos Sei Technol, 67 (2007), 605-15.

13. Vijay S J, Murugan N, "Influence of tool pin profile on the metallurgical and mechanical properties of friction stir welded Al-10 wt.% TiB2 metal matrix composite," Mater Des., 31 (2010), 3585-9.

14. Gopalakrishnan S, Murugan N, "Prediction of tensile strength of friction stir welded aluminum matrix TiCp paniculate reinforced composite," Mater Des., 32 (2011), 462-7.

15. Elangovan K, Balasubramanian V, S. Babu, "Predicting tensile strength of friction stir welded AA6061 aluminum alloy joints by a mathematical model," Mater Des., 30 (2009), 188-93.

16. Sundaram N S, Murugan N, "Tensile behavior of dissimilar friction stir welded joints of aluminum alloys," Mater Des., 31 (2010), 4184-93.

17. J.V.Wood, P. Davies, J.L.F.Kellie, "Properties of reactively cast aluminium-TiB2 alloys," Mater. Sei. Technol., 9 (1993), 833-840.

18. M. Huang et al., "Effect of in situ TiB2 particle reinforcement on the creep resistance of hypoeutectic Al-12Si alloy," J.Alloys Comp., 389 (2005), 275-280.

19. P.L. Schaffer, L. Arnberg, A.K. Dahle, "Segregation of particles and its influence on the morphology of the eutectic silicon phase in Al-7 wt.% Si alloys," Scripta Mater., 54 (2006), 677-682.

20. H. Yi et al., "Effective elastic moduli of Al-Si composites reinforced in situ with TiB2

particles," Scripta Mater., 54 (2006), 1093-1097. 21. G. E. P. Box, W. H. Hunter, J. S. Hunter, "Statistics for Experiments," John Wiley

Publications, New York, (1978). 22. Babu .S et al., "Optimizing friction stir welding parameters to maximize tensile strength

of AA2219 aluminum alloy joints," Metals and Materials, 15 (2009), 321-330.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Slow-Shot High Pressure Die-Casting (SS-HPDC) Process for AE44 Magnesium Single-Cylinder Engine Block with Short-Fiber Reinforcement in

the Bore Bin Hu1, Pan Wang1, Bob R. Powell2, Xiaoqin Zeng3

'General Motors China Science Lab; No. 56 Jinwan Road; Shanghai, 201206, China 2General Motors R&D Center; 30500 Mound Road, Warren, 48090, USA

'Shanghai Jiao Tong University; No. 800 Dongchuan Road; Shanghai, 200240, China

Keywords: Magnesium matrix composites; High pressure die-casting; Process optimization;

Abstract Single-cylinder magnesium engine blocks with short fiber reinforcement in the bore region were prepared by low speed (slow shot), high pressure die casting to infiltrate the fiber preforms and produce sound castings. The magnesium alloy was AE44 and the cylindrical preforms were made with silica-bonded Saffil fibers. The slow shot, high pressure die casting process (SS-HPDC) was intended to simulate squeeze casting, and processing parameters were identified and optimized. These included the pouring temperature, 760°C; the die temperature, 275°C; the preform preheat temperature, 750°C; and the intensification/infiltration pressure, 90MPa. Their relative contributions to the quality of the cast blocks is discussed and suggestions for further optimization are made. The slow-shot, high pressure die casting process is show to be an effective means for casting the blocks, which are free of porosity can be T4 or T6 heat-treated without metal blistering.

1. Introduction Significant benefits in both fuel economy and vehicle handling will be achieved if the mass

reduction potential of magnesium can be extended to the engine. However, the challenges to this are great. It will be necessary to: (1) overcome the low modulus of magnesium to ensure structural stiffness; (2) provide a low-cost, light-weight wear resistance for the cylinder bores; (3) maintain required mechanical properties and creep resistance at high operating temperatures and loads; (4) control or accommodate the thermal expansion mismatch between magnesium and the other engine components; and (5) preventing corrosion either where magnesium comes in contact dissimilar metals or with the engine coolant. The BMW aluminum-magnesium composite engine, which was introduced in 2005, overcame most of the above problems by replacing the working part of the magnesium engine with an aluminum core, around with a magnesium shell was cast. While this engine remains in production, it is nevertheless a complicated and costly approach to enabling some of the benefits of magnesium for mass reduction.

In the present work we consider the use of magnesium matrix composites to address the above challenges of the Mg engine. Magnesium and aluminum metal matrix composites have received considerable attention for aerospace applications l1,2' and aluminum MMCs have been used in a few automotive applications, including the Honda Prelude cylinder bore [3) and Toyota 2ZZ-GE MMCs cylinder engine bore |4'. An Mg-matrix MMC has the potential to address impart the necessary stiffness and strength to the alloy as well as reducing the thermal expansion of the MMC, making it closer to that of the ferrous components of the crankshaft. The reinforcement may also provide wear resistant to the piston rings in the cylinder bores.

The MMC reinforcement phase chosen for this work is Saffil alumina fibers. The Mg alloy for the matrix phase is AE44. AE44 (Mg-4%A1-4%RE) was developed by Hydro Magnesium and has een in automotive production of die Corvette Z06 engine cradle since 2006112'.

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The decision of whether magnesium matrix composites can be an effective replacement material for aluminum MMCs (material for engine components) is complicated. It requires a critical analysis and in-depth understanding of the applications and services for which they are to be used for and their positional, environmental, and adjoining conditions to the corresponding components. Mechanical and tribological properties, and the interfacial characteristics of the fiber reinforced magnesium matrix alloy have be characterized in previous work 15"7'. This work showed that Mg-MMCs have the potential to satisfy the requirements for cylindrical bore materials

In previous literature, most "melt infiltrated" MMCs have been produced by squeezing casting, however, low productivity and infrequent of this process in the nonferrous casing industry are the main limitations to widespread application of Mg-MMCs '*'. In order to make the concept of linerless magnesium engine block easier to be accepted by the casting and auto industry, we reconsidered the high pressure die casting process for Mg-MMC single engine blocks. In the present report, we discuss a potential alternative to squeeze casting, which we term the slow shot, high pressure die-casting process (SS-HPDC) and draw conclusions about its ability to produce single engine blocks containing metal matrix composite bores.

We have made significant changes to the conventional high pressure die casting process to address known concerns. Because the high ingate speed of the melt can erode and destroy the preform, we redesigned the gating system of the die and corresponding parameters of HPDC process. We reduced the shot speed to reduce the ingate speed. We increased the area of the ingates. We also lowered the melt pouring temperature to extent the solidification time for better infiltration and feeding effect to casting. Eventually, the sound castings with well-infiltrated preforms can be obtained simultaneously.

Our approach then, is to determine the potential of SS-HPDC for infiltrating performs and producing high quality, heat treatable Mg cylinder blocks with Saffil fiber-reinforced cylinder bores. In the work reported, we used AE44 magnesium alloy to Saffil alumina fiber performs to prepare the Mg-MMCs single engine blocks. Future work will broaden the range of alloys for optimum processing and properties in the intended applications.

2. Experimental procedure 2.1 Materials

AE44 ingots were obtained from Nanjing Welbow Metals, Ltd., The alloy composition is shown in Table 1. Both as-received ingot composition and post SS-HPDC cast part compositions are shown. The compositions were determined by inductively coupled plasma-atomic emission spectroscopy (ICP-AES) analysis.

Thermal Ceramics de France supplied silica-bonded, Saffil fiber preforms, as either thick or thin wall cylinders (with 1 degree taper angle in ID (inside dimension) for die release), as shown in Fig.l. These preforms were nominally 15 vol. percent solids (fibers). Chemical and physical characteristics of the fibers are shown in Table 2.

Table 1 Composition (by weight) of AE44 alloy*

Composition, % Al RE Mn Bal. Nominal 3.6-4.4 3.6-4.6 0.18-0.50 Mg

Actual 4.0 4.13(1.7Ce,2.1La,0.17Nd,0.16Pr) 0.27 Mg

*: Hydro® Magnesium composition, and actual composition if from SS-HPDC cast parts

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Fig. 1. Preform dimensions and geometries: # 1 preform (thick wall) is shown left; #2 preform (thin wall), right.

Table 2 Saffil fiber properties for MMC preforms

_ ... , ^„,, Density Specified Fiber diameter Composition (wt%) , , £ r-Z , „, , -, v_ v ' (g/cm ) Fiber vol.% (mean), urn „ „ „ 96-97%Al203fiber 3- , , , « ,« , c Saffil fibers A<>,VÂ I- j 3.3-3.5 15 3-5

4%SiQ2 binder

2.2 Single-cylinder engine block for casting trials Tooling for the single-cylinder block for SS-HPDC was based on the XL-3, 124 cc motorcycle engine block. This commercial engine block is produced by aluminum high pressure die-casting around an iron liner, as shown in Fig.2. In our study, the magnesium blocks were produced by SS-HPDC with AE44 infiltrating the cylindrical Saffil fiber preforms and filling the rest of the cavity to form the cylinder block.

Fig. 2 2-D image of the single cylinder engine block

2.3 HPDC equipment and die The SS-HPDC casting trials were performed on an ÜBE NX650C cold chamber, high

pressure die casting machine. This 650 ton machine is a horizontal-horizontal machine; metal flow is horizontal into the cavity and the motion of the ejector die is horizontal. During these trials, the preforms were preheated in a furnace, which was located close to the die. Preheated performs were transferred quickly from the furnace to the die cavity; the die was then closed; molten AE44 was hand ladled into the shot sleeve. The motion of the piston in the shot sleeve drove the metal shot into the die cavity through the gating system, as shown in Fig. 3 and Fig. 4. The thickness of the gate was greater than the sectional thickness of the casting around preform

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for sequential solidification from casting to feeding system, in order to eliminate the porosity inside the casting.

Fig.3 Dimension of perform and position in the casting

2.4 Process control of SS-HPDC process Some key parameters were controlled for process optimizing: pouring temperature of AE44

metal, thickness and type of preform, and the preheat temperature of preform before positioning in the die cavity. Other parameters, such as the cycle time, die temperature, speed and pressure of shot plunger, were kept at the same value for each casting trials. Fig. 5 shows the shot curve of SS-HPDC process. By comparison with conventional high pressure die casting, the shot speed of melt at ingate in SS-HPDC is about 1/4*, the difference as compare to the traditional HPDC process is that after the filling to the cavity, the infiltration process made a stop during the drop of speed line and the rise of pressure line, as shown in Fig. 4.

Fig.4 Typical speed and pressure curve (shot profile) of the SS-HPDC process for Mg-MMCs cylinder engine block castings

3. Results and discussion During the SS-HPDC casting trials, we selected the pouring of 760°C, we obtained sound

castings with filled radiator and better surface quality. Fig. 5 shows a representative die cast magnesium block, which was casted using the following SS-HPDC process parameters of: 760°C pouring temperature, 0.2-0.3/2.5 (m/s) high/low plunger speed, 275°C die temperature, and 90 MPa intensification pressure (the speed, die temperature and pressure were very popular for conventional HPDC process for better acceptance by industry, the high pouring temperature and gate design of die are the biggest difference between SS-HPDC and conventional HPDC process).In these experiments we determined and other process parameter, that of the temperature of the preform. When casting was done with preforms preheated to 450°C, the

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block portion of the magnesium casting was sound, but there was no observed infiltration of AE44 into the preforms. The preform can't be infiltrated well until the preheat temperature was gradually increased to 750°C.

Fig.5 SS-HPDC AE44 Magnesium block with well-infiltrated perform

Both thick wall and thin wall preforms were infiltrated by SS-HPDC process using 750°C preheat temperature (other parameters were as note before). Fig. 6 shows cross-section images of the SS-HPDC casting with different preforms. The regions of the preforms at either end are slightly deformed, and some shrinkage pores were found in the casting with the thick preform (#1). Because of the same OD (outside diameter) of two preforms (the thicker preform have 2mm smaller ID), so the casting with thicker perform have more volume of Mg alloys in the die cavity. And the feeding efficiency of gating system was limited in the casting with thick preform, more porosities were found in the casting with thick preforms.

Fig. 6 Sectional image of the casting: (a) with thick perform, # 1 ; (b) with thin perform, #2

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Fig. 7 shows the microstructure of an HPDC-infiltrated thin perform (#2). There are two grain sizes in the casting, ~30um and ~5um finer grain size, which probably formed in the sleeve/gating and cavity, respectively, as shown in Fig. 7(b). Some inclusions and oxidation products can be found in the interface between the casting and the preform, as shown in Fig. 7 (c), which probably formed during the filling of Mg melt in the sleeve and cavity. The surface of the preform appears to have functioned like a filter. About 0.5mm thick of un-reinforced region was found in the region which contact with the core the die, as shown in Fig. 7(d).This ascribed to the low temperature of the perform, which contact with the die, limited the infiltration of the melt. Fig. 8 shows the comparison of microstructure between SS-HPDC process and squeeze casting process (SQCI?1), grain size of SS-HPDC casting are much finer than SQC process (very popular for preparing metal matrix composite), both in composite region and unreinforced matrix alloy region.

Fig.8 SEM microstructure of the casting: (a) Region in casting with preform; (b) Matrix alloy without fiber(c) Interface between matrix and composite; (d) Edge of composite

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Fig.8 Optical microstructure of the casting by SQC and SS-HPDC process: (a) composite by SQC; (b) matrix by SQC; (c) composite by SS-HPDC; (d) matrix by SQC

Above discussion implied that the sound single engine block casting with well-infiltrated preform can be obtained by SS-HPDC process. The grain size of SS-HPDC was much finer than SQC process, which means the mechanical property can be probably higher than SQC process. However, traditional HPDC process have more gas entrapment during filling of the die cavity, so the traditional HPDC castings had much lower strength than SQC castings even if the grain size is much finer. The SS-HPDC was expected to ultra-low gas entrapment (can be heat-treated for better mechanical properties) due to much lower filling speed by large gating system.

In order to test the porosity containing of the Mg-MMCs block produced by SS-HPDC process, block with #2 preform was selected to carry out the T4 heat-treatment (too much porosity were found in the blocks with # 1 preform, as shown in Fig. 6a). Fig. 9 shows the T4 heat-treated single cylinder engine block produced by SS-HPDC process (contained #2 preform, 450°CX6h+air cooling), no surface blister was found in casting after heat-treatment. The SS-HPDC process can produce the porosity-free castings because no air-entrapment of the melt by the low ingate filling speed of melt.

Fig.9 Surface condition of the heat-treated SS-HPDC Mg-MMCs block (with #2 preform, 450°CX6h+air cooling)

Conclusions The slow shot, high pressure die casting process (SS-HPDC) effectively infiltrated the fiber

perform and yielded a sound Mg cylinder block around the preform. SS-HPDC may be a viable alternative to squeeze casting. Key parameters were optimized to obtain the sound castings and well-infiltrated preforms simultaneously.

The optimized process parameters were: high/low plunger speed, 0.2-0.3/2.5 (m/s); pouring temperature, 760°C; die temperature, 275°C; highest solidification pressure, 90MPa; and preform preheat temperature, 750°C.

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3. The thickness of the preform should be as thin as is consistent with subsequent machining tolerances in order to ensure good infiltration (thin for ease of infiltration and thick enough to have sufficient functional material at the bore ID after finish machining). The thickness of the ingate should be thicker than the sectional thickness of the casting around perform for sequential solidification from casting to gating system.

4. Heat treatability of castings can be obtained by SS-HPDC process without blister after solution treatment (450°CX6h). Microstructure of SS-HPDC casting is much finer than that made by squeeze casting process (both composite region and unreinforced casting region). Heat-treatable alloy will be selected for the matrix alloy in the future investigation, for better mechanical properties and wear-resistant properties.

Acknowledgements The authors would like to thank LAF engineers Xiaoyong Chen and Wenchun Cai for their kindly and considerable assistance during the preparation of the castings.

Refernces 1. K.U. Kainer, ed., Metal Matrix Composites, Wiley-VCH, Weinheim, Germany, 2006. ( K.U.

Kainer, Metal matrix composites. Custom-made materials for automotive and aerospace engineering, Wiley-VCH, Weinheim).

2. N. Chawla and K.K. Chawla, Metal Matrix Composites, Springer, New York, USA, 2006. 3. N. Chawla and K.K. Chawla, "Metal Matrix Composites in Ground Transportation," JOM,

58, 2006, 67-70. 4. Toshihiro Takami, et al., "MMC All Aluminum Cylinder Block for High Power Si Engines",

SAE Technical paper, No. 2000011231, 2000. 5. B. Hu, L. Peng, B. Powell, and A. Sachdev, Mechanical Properties and Microstructural

Analysis of AXJ530 Magnesium Alloy Reinforced with Alumina Fibers, Magnesium Technology (The Minerals, Metals & Materials Society, TMS), 2010, 473-480

6. B. Hu, L. Peng, B. Powell, M. Lukitsch and A. Sachdev, Dry Sliding Wear Behavior of AE44 Magnesium Alloy Reinforced with Saffil Alumina Fibers, Magnesium Technology (The Minerals, Metals & Materials Society, TMS), 2010,297-303

7. B. Hu, L. Peng, B. R. Powell, M. P. Balough, R. C. Kubic, A. K. Sachdev, Interfacial and Fracture Behavior of Short-Fibers reinforced AE44 Based Magnesium Matrix Composites, Journal of Alloys and Compounds, 504,2010, 527-5345

8. H. E. Friedrich, B. L. Mordike, Magnesium Technology, metallurgy, design data, applications. Springer, 2005. (6.3 Magnesium matrix composite, Page 324)

9. K.E. Nelson, "Magnesium Die Casting Alloys" (Paper presented at the 6th SDCE International Die Casting Congress, Cleveland, Ohio, 1970), paper no. 13.

10. Y.D. Huang et al. "Evolution of microstructure and hardness of AE42 alloy after heat treatments", Journal of Alloys and Compounds, 463 2008, pp.238-245.

11. P. Bakke, H. Westengen, The role of rare earth elements in structure and properties control of magnesium die casting alloys, in: N.R. Neelamegham, H.I. Kaplan, B.R. Powell(Eds.), Proceedings of the Magnesium Technology, TMS, Warrendale,2005, pp. 291-296.

12. N. Li, R. Osbome, B. Cox, D. Penrod, The USCAR structural cast magnesium development project, in: N.R. Neelamegham, H.I. Kaplan, B.R. Powell (Eds.), Proceedings of the Magnesium Technology, TMS, Warrendale, 2005, p. 535.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

COMPRESSIVE PROPERTIES OF A1-B4C COMPOSITES OVER THE TEMPERATURE RANGE OF 25 - 50ffC

Srinu Gangolu1, A. Gourav Rao2, N. Prabhu1, V.P. Deshmukh2, B. P. Kashyap1

'Department of Metallurgical Engineering and Materials Science, Indian Institute of Technology Bombay, Mumbai, 400076, India

2Naval Materials Research Laboratory, Shil Badlapur Road, Addl. Ambemath (E) Dist, Thane 421506, India

K e y w o r d s : AI-B4C composite, Boron carbide, Stress-strain curves, Temperature effect, Grain size

Abstract

Aluminum-Boron carbide (AI-B4C) composites containing Al-0, 5, 10 and 15 wt% B4C were produced by flux assisted reaction method at 900°C. In order to study the effects of temperature on flow property and microstructural evolution, compression tests were done over the temperature range of 25° to 500°C at an initial strain rate of 10"2 s"' and the microstructures were examined in Scanning Electron Microscope. The stress-strain curves exhibit an increase in flow stress with the B4C content, except that the same decreased for Al-15 wt% B4C composite below that of Al-10 wt% B4C composite. The flow stress was found to decrease with the increase in test temperature above ~200°C, but it exhibited nearly athermal behavior below this critical temperature. The flow property of the composite is discussed in terms of the flow of the matrix and particulate, and the concurrent microstructural evolution that occurs during deformation.

Introduction

Aluminum with good tensile ductility is an attractive light metal for applications in various industries but it does not possess adequate strength for structural applications. This has necessitated development of various Al-based alloys and composites to exploit the strengthening effects of alloying elements and reinforcements. Metal matrix composites are attractive materials for automobile and aerospace applications due to their light weight and high strength [1], Aluminum is commonly reinforced with many ceramic particles such as SiC, AI2O3, A1N, TiC and B4C. While Al with SiC and AI2O3 particulates have been widely investigated, there exists only limited studies on AI-B4C composites. Recently, AI-B4C composites have received new attention in particulate metal matrix composites (PMMC) due their high temperature mechanical properties, low density and high hardness, chemical resistance, and neutron absorption property [2,3],

The mechanical behavior of MMCs depends on the interface formed due to chemical reactions between matrix and particle. Viala et al [4] reported the chemical reactivity between Boron carbide and aluminum. In this system, the different phases, viz AIB2, AI3BC, AI4C3 and AI3B48C2, are formed depending on the heat treatment temperature, which can have influence in the property of AI-B4C composites. In the past, microstructural evaluation of composite materials was generally done using optical and scanning electron microscopes but, owing to the high interfacial area in composites, through particle/matrix interface and grain boundaries by grain refinement, it used to be difficult to study effectively. Recently, electron backscatter

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diffraction (EBSD) technique is used as a powerful characterizing tool for better understanding [5, 6], A few investigations have been reported on the deformation of AI-B4C composites over wide ranges of temperatures till the date [3, 7, 8]; Onoro et al. [3] investigated the high temperature behavior of Al 6061 and 7015 composites reinforced with B4C particles from room temperature to 500°C. However, no such study seems to be present on AI-B4C composites based on commercial aluminum. A study on flow properties, especially in compression, could be useful for hot working of composites for better properties and integrity over the as-cast composites. Hence, it is important to investigate the high temperature deformation behavior of commercial aluminum reinforced with B4C particles to understand effective strengthening of AMCs at various volume fractions and temperatures. Therefore, in the present work, compressive properties of AI-B4C composites were investigated over a wide range of temperatures, ranging from room temperature to 500°C, at an initial strain rate of 10"2 s"1 and the microstructural characterization was made by using analytical microscopy techniques, including EBSD, to evaluate the distribution of B4C reinforcement and the grain size of composites.

Experimental Procedure

In this investigation, commercial aluminum was used as matrix and boron carbide particles of average size 21 um as reinforcement. Aluminum-boron carbide composites of 0, 5, 10, and 15 wt% were produced by flux assisted reaction method [9,10] at 900°C. Potassium hexa-flourotitanate flux (KaTiFö), in the proportion of 10% of B4C content, was used for enhancing die wettability of ceramic particles in liquid aluminum. A 2 kg of aluminum was initially melted in graphite crucible at 900°C and then the mixture of boron carbide with flux was added to liquid aluminum at this temperature. For microstructural studies, the composites were initially polished with 220, 400, 600, 800, 1000 and 1200 grade SiC papers, which was followed by polishing with alumina slurry and diamond paste of sizes 6, 0.5 and 0.25 urn sequentially. Keller's reagent consisting of HNC>3:HC1:HF:H20 in a proportion of 5:3:2:190 was used as an etchant for 60 seconds. The samples were also polished with SiC>2 colloidal suspension then subjected to electropolishing for 17 S of dwell time at 13 V for electron backscatter diffraction studies.

The compression samples of dimensions 9 mm length and 6 mm diameter were machined from the as-cast AI-B4C composites. Compression tests were done from room temperature to 500°C at a strain rate of 10"2 s'1 using Zwick- Roell Amsler Universal Testing Machine of 100 kN capacity. The heating and soaking time elapsed prior to compression test was 60 min. All the samples were compressed to 2 mm final thickness. The microstructures of the deformed composites were examined in loading and transverse directions by Hitachi S-3400 model Scanning electron microscope. EBSD was done over the electropolished area of 85 um x 150 um at a step size of 0.1 (im.

Results and Discussion

Preliminary Characterization

Figure 1(a) (b) (c) and (d) shows the scanning electron microscope images of Al, Al+5% B4C, Al+10% B4C and Al+15% B4C composites, respectively. It is clear that particle distribution is uniform these composites but the interparticle distance decreases as the fraction of reinforcement increases. Although the particle distribution seems to be homogenous, further examination at higher magnification and its compositional analysis by Energy Dispersive Spectroscopy (EDS) reveal a variation in composition from the particle/matrix interface towards the interior into the

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matrix. Figure 2 (a) shows an EDS of Al-10% B4C composite with the selected reference line for analysis marked. It was observed that potassium hexa-flourotitanate (K^TiFö) flux added for better wettability formed clusters near the particle instead of enhancing the wettability of B4C particle by forming an interfacial layer. Figure 2 (b) shows an EDS analysis of Al-10% B4C composite starting from AI-B4C interface marked in Fig. 2(a), which reveals the variations in concentration of aluminum and titanium within the micrograph. Toptan et al. [11] reported the titanium layer of 180-80 nm thickness around the particle, but Fig .2 shows the flux to be distributed to few microns instead of remaining at the particle /matrix interface. However, similar analysis in the other composites, viz. Al-5% B4C and Al-15% B4C, did not reveal such a compositional variation from the interface.

Figure 1. SEM images of as-cast (a) Al; (b) Al+5% B4C; (c) Al+10% B4C and (d) Al+15% B4C composites

Figure 2. EDS analysis of Al+10% B4C composite (a) with the variation in Al and Ti contents (b) along the line marked with arrow in (a)

Effect of Reinforcement on Compressive Property

True plastic stress- true plastic strain curves of Al, Al+5% B4C, Al+10% B4C and Al+15% B4C typically at two test temperatures of 25° and 400°C are presented in Fig. 3. As the weight fraction of reinforcement increases, the flow curve moves upwards except for 15% reinforcement, which shows the lower flow stress than that of Al+10% B4C composite. This kind of anomalous flow behavior was also reported by Mazahery and Shabani [12] in squeeze cast A356-B4C composite for the same levels of reinforcements. The reason for this anomaly is given in terms of higher internal stress built-up in the Al-15% B4C composite as compared to the Al-10% B4C composite. The nature of flow curves of the composites at all the test temperatures

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investigated were nearly the same. It is clear from room temperature plot in Fig. 3(a) that the slope of stress-strain curves (work hardening rate) corresponding to the Al, Al-5% and Al-15% composites were not changed with reinforcement, but the curve for the Al-10% B4C composite shows slightly greater slope. However, at 400°C test temperature, the stress-strain curves exhibit similar kind of slope at early part of straining but, at larger strain (true stain 0.6), the stress-strain curve for Al-5% B4C composite is noted to cross over that for the Al-15% B4C composite. The gap between two stress-strain curves represents the reinforcement effect irrespective of temperatures at which the sample is tested. The increase in flow stress with the increasing reinforcement is attributed to the increasing pinning effect by B4C particles to dislocation motion.

Figure 3. True stress-true strain curves of AI+B4C composites at (a) Room Temperature and (b) 400°C

Effect of Temperature on Compressive Property

The variations in flow stress of all AI+B4C composites, as a function of temperature, are shown in Fig. 4 typically at a true strain of 0.10. The nature of variation in flow stress as a function of temperature suggests two distinct regions, (i) There occurs no change in flow stress with the increase in temperature up to -200 in most of the composites. This behavior is similar to the athermal behavior that generally occurs at intermediate temperatures of about 0.3-0.7 Tm, Tm being absolute melting point. However, the close examination of the plot shows that the limiting temperatures of this athermal behavior are changed with the fraction of the reinforcement. For example, Al-0% B4C shows athermal behavior up to 300°C whereas Al-10% B4C suggests the existence of low temperature thermal activation for deformation, (ii) The flow stress decreases rapidly with the increase in temperature. This behavior is typical of high temperature deformation, which is facilitated by diffusion and grain boundary sliding. The transition temperature between the athermal behavior and the high temperature behavior is reported [13] to decrease with the increase in flow stress with strain and the microstructural changes that occur during deformation. For example, in type 316L stainless steel, the work hardening and the associated substructure development with increasing strain lowers this transition temperature, requiring less thermal activation for flow due to the enhanced contribution of stress itself by work hardening.

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Figure 4. Variation in flow stress with temperature at true strain of 0.1

Microstructural Evolution

Microstructures of the Al+5% B4C and Al+15% B4C composites, undergone compression tests at room temperature and 400°C at strain rate of 10"2 s"1, were examined in the loading and transverse directions. Presented in Figure 5(a), (b) are microstructures of Al+5% B4C composite and 5(c), (d) are microstructures of Al+15% B4C composites deformed at room temperature and 400°C, respectively. Here all the micrographs correspond to the transverse sections of the deformed samples to engineering strains shown as inset in the micrographs. It is clear from comparison of these micrographs with the initial microstructures shown in Fig. 1(b), (d) that B4C particles get preferentially aligned in transverse direction upon deformation, which is also the direction of material flow in compression. Where applicable based on alignment of particles, the transverse directions (TD) in the micrographs are marked by arrows. As the test temperature increases the inter-particle distance in the transverse direction is noted to decrease. Since the movement of particles is not the same as that of matrix, the B4C hard particles tend to preferentially get aligned in the direction of material flow normal to the compressive loading direction, especially when the test temperature becomes high.

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Figure 5. Microstructures of composites in transverse sections upon deformation: (a) Al+5% B4C at room temperature, (b) Al+5% B4C at 400°C, (c) Al+15% B4C at room temperature and (d)

Al+15% B4C at 400°C. (Arrow indicates Transverse Direction - TD)

Electron Backscatter Diffraction

Figure 6 shows EBSD maps for Aluminium and Al+5%, Al+10% B4C composites all deformed at room temperature. These maps were indexed by considering aluminium. Due to higher magnification and conventional polishing technique employed the quality of EBSD maps could not bring out the important features in details. However, the average grain size and the misorientation angles obtained in Al and its composites here are listed in Table 1.

Figure 6. EBSD of (a) Aluminum, (b) Al+5% B4C and (c) Al+10% B4C composites, all subjected to compression tests at room temperature.

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Table 1. Grain size and misorientation angles of AI-B4C composites as obtained from EBSD

Aluminum Al + 5% B4C Al +10% B4C

Grain size (um) 32.6 11.8 2.6

Misorientation angle 30.9 35.2 39.1

It is seen that as the reinforcement increases the grain size reduces. The flow behavior of composites in Fig. 3 and Fig. 4 follows the Hall-Petch type relationship, with the smaller grain size showing higher flow stress. Fig 6(a) shows the elongated grains of aluminum in transverse direction whereas the composites still show equiaxied grains after compression tests. The average misorientation angle is also higher for Al+10% B4C composite. The flow properties of the composites in the present work are seen to be the results of not only the fraction of B4C but also its effect on grain size and the change in microstructure that occurs in the course of deformation.

Conclusions

Compression tests on Al and its composites containing 5, 10, and 15 wt% B4C at strain rate 10"2

s"1 and the microstructural examination of deformed composites leads to the following conclusions. Flow stress increases with increase in B4C content and decrease in test temperature. However, there occur some deviations from this trend towards the lower test temperature and the higher B4C level of 15 wt%. Some of these anomalies can be understood by the nature initial distributions of B4C particles and the concomitant changes in grain size, misorientation angle and the particle redistributions in the course of deformation.

References

1. M.K. Surappa, "Aluminium Matrix Composites: Challenges and Opportunities," Sadhana, 28 (2003), 319-334.

2. R.M. Mohanty and K. Balasubramanian, "Boron Rich Boron Carbide; An Emerging High Performance Material," Key Eng. Mater., 395 (2009), 125-142.

3. J. Onoro, M.D. Salvador, and L.E.O. Cambronero, "High-temperature mechanical properties of aluminium alloys reinforced with boron carbide particles," Mater. Sei. Eng., A, 499 (2009), 421-426.

4. J.C. Viala, J. Bouix, G. Gonzalez, and C. Esnouf, "Chemical reactivity of aluminium with boron carbide," J. Mater. Set, 32 (1997), 4559-4573.

5. J. Guo, S. Amira, P. Gougeon, and X.G. Chen, "Effect of the surface preparation techniques on the EBSD analysis of a friction stir welded AAHOO-B4C metal matrix composite," Mater. Charact, 62 (2011), 865-877.

6. X.G. Chen M.D. Silva, P. Gougeon, and L.S. Georges, "Microstructure and mechanical properties of friction stir welded AA6063-B4C metal matrix composites," Mater. Sei. Eng., A, 518(2009), 174-184.

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7. C. Nie, J. Gu, J. Liu, and D. Zhanget, "Deformation and fracture behaviour of 7039 Al reinforced with B4C particles at elevated temperature," Key Eng. Mater., 351 (2007), 65-69.

8. W. Kai, J.M. Yang, and W.C. Harrigan, "Mechanical behaviour of B4C particulate-Reinforced7091 aluminium composite," Scr. Metall., 23 (1989), 1277-1280.

9. A.G. Rao, M. Mohape, V.A. Katkar, D.S. Gowtam, V.P. Deshmukh and A.K. Shah, "Fabrication and Characterization of Aluminium (6061)-Boron-Carbide Functionally Gradient Material," Mater. Manuf. Processes, 25 (2010), 572-576.

10. A.R. Kennedy and B. Brampton, "The reactive wetting and incorporation of B4C particles into molten aluminium," Scr. Mater., 44 (7) (2001), 1077-1082.

11. F. Toptan, A. Kilicarslan, A. Karaaslan, M. Cigdem, and I. Kerti, "Processing and microstructural characterisation of AA 1070 and AA 6063 matrix B4CP reinforced composites," Mater. Des., 31 (2010), 87-91.

12. A. Mazahery and M.O. Shabani, "Mechanical Properties of Squeeze-Cast A356 Composites Reinforced with B4C Particulates," J. Mater. Eng. Perform., (2011), 1-6.

13. B.P. Kashyap, K. McTaggart, and K. Tangri, "Study on the substructure evolution and flow behaviour in type 316L stainless steel over the temperature range 21-900^," Philos. Mag. /4, 57(1) (1988), 97-114.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

ALUMINUM METAL MATRIX COMPOSITE VIA DIRECT METAL LASER DEPOSITION: PROCESSING AND CHARACTERIZATION

B. L. Waldera and S. J. Kalita

Advanced Engineered Materials Center School of Engineering and Mines, University of North Dakota

Grand Forks, North Dakota, United States

Keywords: aluminum metal matrix composite, direct metal laser deposition, nahoindentation

Abstract

Aluminum metal matrix composites (Al-MMCs) have shown favorable material characteristics for aerospace applications such as airframes, reinforcement materials and joining elements. Similarly, carbide Al-MMC coatings can greatly improve surface performance of aluminum components. In this research, such coatings were developed from a powder blend of pure aluminum, chromium carbide and tungsten carbide nickel alloy, premixed at 2:1:1, on AA 7075 plates through direct metal laser deposition. Microstructure of the carbide Al-MMC was studied by optical and scanning electron microscopy. The hardness and the reduced Young's modulus (Er) were assessed through depth-sensing instrumented indentation using a Hysitron-TI950 Nanomechanical Tester. The carbide Al-MMCs demonstrated good interfacial bonding and improved modulus and hardness. A maximum hardness and reduced modulus of 2.47 GPa and 87.43 GPa, respectively, were recorded for the composite. Hardness and reduced modulus of AA 7075 substrate were 1.45 GPa and 70.63 GPa, respectively. This research presents our recent findings.

Introduction

Aluminum alloys have been used extensively in aerospace, automotive and marine applications due to their high specific strength, low density, good electric and thermal conductivity and good workability. However, improvement of surface properties, such as hardness and wear, is an area of extant research and development. Al-MMC coatings have shown a propensity for greatly improving the surface performance of aluminum components. Al-MMCs offer advantages over conventional alloys when the property profile of a standard material is no longer sufficient for a specific application [1]. Novel materials may be custom made with advanced properties depending on the requirements of the application. The possibility of combining metals, ceramics, and non-metals yields virtually endless combinations of materials with a broad spectrum of properties. Carbides such as SiC, TiC and WC are often used to form ceramic-metal composites on the surface of conventional alloys in order to enhance wear resistance [3-5], WC-based composites are known to have excellent wear resistance-strength combination [6].These types of materials become especially advantageous when there is a reasonable cost-performance relationship in the component production.

Direct Metal Laser Deposition (DMLD) is a fabrication technique for building components in a layer by layer fashion. In this process, a laser beam is used to melt powder onto the surface of a substrate creating a metallurgical bond with the substrate after rapid solidification of the melt pool. With powder injection DMLD, metal powder is injected directly into the laser beam by an inert gas flow. The energy in the laser beam melts the in-flight particles allowing a melt pool to form with subsequent solidification of material on the substrate. Since

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DMLD offers low heat input, MMCs can be fabricated, carbides intact, with minimal dissolution and melt only the matrix material [2], Slight dilution of the substrate into the melt region allows an excellent metallurgical bond to develop. Also, laser processing offers advantages over other fabrication methods in terms of processing as well as the resulting microstructures achieved.

Studies on improving mechanical characteristics of aluminum alloys through laser surface treatment are abundant. Many studies aimed at forming intermetallic aluminum compounds to improve properties. Also, the development of MMCs with hard reinforcement particles distributed in a softer metal matrix has been evaluated with reinforcement particles like silicon carbide and titanium carbide [7, 8]. The current research demonstrated fabrication of Al-MMCs and investigated the microstructure and nanomechanical properties of DMLD Al-MMC coating on AA 7075 from a powder blend of pure aluminum, chromium carbide and tungsten carbide nickel alloy, premixed at 2:1:1.

Experimental Procedure Laser Processing

DMLD was performed using an IPG 1KW fiber coupled diode laser, which generates a 1070-1080 nm wavelength with a continuous wave mode. The power, cooling and control for the laser system were provided by peripheral units. The CNC 5-axis control provided precise movement of the specimen relative to the laser beam. The beam was focused on the surface of the specimen, using an objective lens integrated into Precetec YC50 cladding head. YC50 assembly with 100 mm FL Collimator and 200 mm focal length produces a theoretical 800 micron (0.032") spot with 400 micron fiber. The powder composition consisted of 50 wt% aluminum powder (Metco 54NS), 25 wt% Chromium Carbide (Metco 70C-NS) and 25 wt% Tungsten Carbide Nickel Superalloy (SM 5803). A 50 liter/min flow of argon gas through the surrounding annulus shrouded the process from oxidation. The specimens were ground to 600 grit and cleaned with isopropyl alcohol prior to DMLD. The continuous wave mode was employed using laser power, speed and strategy as described in Table 1 and Figure 1.

Table 1. Laser Processing Parameters Name

AA7075 600W600 600W800 600W1000 700W1300 700W1600

Laser Power (W)

-600 600 600 700 700

Travel Speed (mm/min)

-600 800 1000 1300 1600

Laser Specific Energy (J/mm2)

-75.0 56.3 45.0 40.4 32.8

Hatch Spacing (mm)

-0.6 0.6 0.6 0.6 0.6

Figure 1. DMLD Strategy and the Resulting Thick Al-MMC composite Coating

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Microstructural Analysis Specimens were mounted in resin blocks and prepared through standard metallurgical

procedures. Before microstructural examination, the samples were ground using SiC grinding paper with 1200 grit size then polished until 0.5 pun Si02 suspension. The polished specimens were chemically treated with vilella's etchant (1 g picric acid, 5 ml HCL, 100 ml ethanol) for approximately 30s. After etching, the specimens were cleaned by demineralized water and isopropyl alcohol. Specimens were characterized using a Zeiss confocal microscope fitted with AxioCam MRc 5 for optical microscopy examination. Additional microscopy was performed using a Hitachi S-3400N Scanning Electron Microscope.

Nanomechanical Testing Nanoindentation experiments were conducted with a diamond Berkovich tip. The tip area

function was calibrated on a Fused Quartz sample with reduced modulus (Er) = 69.6 GPa and hardness (H) = 9.25 GPa. Friction stir weld samples were prepared by grinding on silicon carbide abrasive paper to 1200 grit followed by polishing with a 6 um diamond suspension, 1 um diamond suspension and finally 0.05 um SiC>2 suspension. Thirty-six indentations were made on each sample with a peak force of 6000 uN. The loading rate was 1200 uN/s with 5 s loading and unloading periods. Unload occurred 10 s after the peak force was reached.

The hardness and reduced modulus were determined using the multicurve analysis option in TriboScan® version 8.1.1 (Hystron Ine, Eden Prärie, MN). This process for determining hardness and reduced modulus was in accordance with the standard method for indentation data analysis outlined by Oliver and Pharr [9], First, the unloading curve is fitted with a power law relation. The Upper Fit % and Lower Fit % parameters were 95% and 20%, respectively. This indicates the portion of the unloading curve to fit the power law relation (1)

P = A(h-hf)m (1)

where P is the indentation load, h is penetration depth, hf is the residual displacement after unloading, and A and m are determined by the fitting procedure. Hardness is calculated with equation (4) where A(hc) is the area function. Reduced modulus is calculated with equation (5).

Results and Discussion Microstructure

Figure 2 presents the optical microscopy as observed with the Zeis confocal microscope. The thickest clad region can be observed in the 600W800 specimen. Furthermore, the 600W600 through 700W1300 specimen exhibits evidence of a carbide layer on the top surface of the clad zone (CZ). On 700W1600 there is sufficient dilution of the carbide layer resulting in no carbide accumulation on the surface of the sample. Carbide particles are instead mixed throughout the CZ. This is further confirmed in SEM analysis of Figure 3. The swirling effect observed in the CZ is due to high temperature gradients within the melt pool which causes the liquid phase to mix rapidly as a result of convection forces developed by in the melt pool.

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Figure 2. Optical micrographs of Al-MMC composite coating, processed using different laser parameters.

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Figure 3. Scanning electron micrographs of Al-MMC composite coating, processed using different laser

Nanoindentation Figure 4a presents reduced modulus as a function of displacement for the tested

specimens. The data collected behaved in a linear manner from sample to sample which is predicted by the nature of nanoindentation as outlined by Oliver and Pharr. Figure 5a again confirms that the reduced modulus was increased in the 700W1600 sample while there is no significant difference in hardness for the other specimens. Similarly, nanohardness values behaved in an analogous fashion. Displacement however, is inversely proportional to nanohardness as outlined by equation 4.

A maximum hardness and reduced modulus of 2.47 GPa and 87.43 GPa, respectively, were recorded for the composite in sample 700W1600. Hardness and reduced modulus of AA 7075 substrate were 1.45 GPa and 70.63 GPa, respectively. Samples 600W600, 600W800, 600W1000, and 700W1300 indicated no improvement in hardness as seen in figure 5. This however is not representative of a homogeneous clad zone. This is primarily indicative of the matrix material and not representative of carbide particles dispersed throughout the melt region.

In DMLD, the laser beam melts the metal matrix material but does not completely melt the ceramic carbide particles which are injected into the melt pool. This leaves carbide particles scattered throughout the clad region as observed in particle reinforced composites. The values measured in four of the five DMLD specimens appear to be representative of the matrix material.

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The fifth 700W1600 shows evidence of carbide particles in the melt region. This is evident by observing the increased average hardness and reduced modulus (see figure 4 and figure 5) but also by the standard deviation of the hardness values. Regions with greater amounts of carbide particles will have higher hardness and reduced modulus values than regions where only aluminum matrix material exists. Thus nanoindentation appears to be a less desirable choice for determining the hardness throughout the melt region as an entity because of its small scan area. It is however, an excellent method for determining the hardness of individual carbide particles however, motion control allowing this is not available on the Hysitron-TI950 Nanomechanical Tester.

Figure 4. Reduced modulus (A) and nanohardness (B) of Al-MMC processed with 700W laser power with a scanning speed of 1600 mm/min as a function of displacement.

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Figure 5. Statistics for (A) Reduced Modulus and (B) Hardness vs. Displacement

Conclusions The microstructure and nanomechanical properties of DMLD coatings of a powder blend

of pure aluminum, chromium carbide and tungsten carbide nickel alloy, premixed at 2:1:1, deposited on AA 7075 were investigated. Specimens indicated a swirling effect in the microstructure as a result of high temperature gradients within the melt pool. This caused the liquid phase to mix rapidly as a result of convection forces developed in the melt pool. This movement of molten material allowed the accumulation of carbide particles on the surface of four out of five DMLD specimens. One specimen exhibited sufficient dilution of the carbide layer resulting in no carbide accumulation on the surface of the sample.

Nanoindentation indicated a possible increase in hardness of the melt region as a result of carbide particles mixing with the aluminum matrix. Nanoindentation however, prevented large areas to be tested which likely resulted in direct observation of the matrix material in four of the 5 specimens. Microhardness may be a better approach in future iterations.

The present study demonstrates the capability to produce MMCs of this specific carbide formulation with particles observed in the melt pool promoting increased hardness and reduced elastic modulus for future applications in advanced material science.

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Acknowledgements

The authors are grateful of the following organizations for their support and the success of this program: US Army Armament Research, Development and Engineering Center (ARDEC) and Benêt Laboratories. US Army Benêt Laboratories and die Advanced Engineered Materials (AEM) Center of the University of North Dakota are working under a joint program entitled 'Lightweight Reliable Materials for Military Systems' (Agreement No: W15QKN-11-2-0002). Special thanks to Benet Laboratories Cooperative agreement Manger (CAM) Mr. Kevin Miner, and Alternate Cooperative agreement Manger (ACAM) Mr. Fang Yee for their suggestions and help in managing and executing this research program.

References

[1] Karl U. Kainer, "Metal Matrix Composites." Wiley-VCH. 2006 (95) [2] J. Nurminen et al. / Int. Journal of Refractory Metals & Hard materials. 27 (2009) 472-478 [3] H.C. Man, S. Zhang, F.T. Cheng and T.M. Yue. Scripta Materiala. 44 (2001) 2801-2807 [4] M. Cadenas, R. Vijande, H.T. Montes, and J.M. Sierra, Wear. 212, 244 (1997) [5] P. Wang, Y. Yang, G. Ding, J. Qu, and H. Shao, Wear. 209, 96 (1997) [6] J. Sampedro, et al. Physics Procedia. 12 (2011) 312-322 [7] J.M. Pelletier et al. J. Phys. IV 1 (1991) 61-63 [8] L. Dubourg et al./ Wear 258 (2005) 1745-1754 [9] W.C. Oliver, G.M. Pharr,./. Mater. Res. 7 (1992) 1564.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

DAMAGE EVOLUTION MODEL FOR HYBRID METAL MATRIX COMPOSITES

Jessica A. Dibelka1, Scott W. Case1

'Department of Engineering Science & Mechanics Virginia Polytechnic Institute and State University

223 Norris Hall; Blacksburg, VA, 24061, USA

Keywords: Metal matrix composite, tensile strength, progressive failure

Abstract

In this study, two types of alumina fiber reinforcement, continuous Nextel™ fabric and discontinuous Saffil paper, were used to produce aluminum matrix composites. One of the composites created was a hybrid composite consisting of both reinforcements. This hybrid composite displayed high strength similar to the Nextel™ fabric-reinforced composite while maintaining a reasonable strain to failure typical of the Saffil paper-reinforced composite. Additionally, the hybrid composite exhibited progressive rather than catastrophic tensile failure. When the maximum strength of the composite was reached, the stress dropped 10% and was sustained for a considerable amount of strain before ultimate failure. A modeling framework is presented to describe the progressive damage and ultimate failure of these composites. This framework combines fiber/matrix-scale models for stiffness and strength with layer-level models for stress redistribution. Excellent predictions are obtained for initial failures, with reasonable predictions obtained for subsequent damage events.

Introduction

Alumina reinforced aluminum composites have an increased specific strength and specific stiffness [1] but a decreased fracture toughness compared to aluminum alloys. Characterization of hybrid reinforced aluminum composites with continuous and discontinuous alumina fiber reinforcements have become a priority due to the possibility of developing materials with an enhanced toughness. Tension testing of a hybrid composite reinforced with alumina fibers in two different geometries—a woven Nextel™ fabric and a discontinuous randomly oriented Saffil paper—revealed an interesting non-catastrophic failure mechanism. During tension testing it was apparent that the outer Nextel™ fabric layer cracked causing a 10% reduction in stress while the Saffil paper layer remained intact and continued to carry loading for an additional 0.05% strain. Micromechanics based modeling was used to further the understanding of how the non-catastrophic failure mechanism occurred and to predict the stress strain behavior of the composite. The hybrid composites' stress strain behavior was predicted using three methods. First, the experimental tensile stress-strain results for the two base composites (those reinforced with either Nextel™ fabric or Saffil paper) were used in a rule-of-mixtures approach to predict the stress-strain response of the hybrid composite. Then, Xia and Curtin's [2] continuous fiber reinforced composite strength model was used to predict the cracking strength of the hybrid composite through a damage evolution scheme. Finally, Hashin's [3] variational calculus based stress analysis for composites containing cracks was used to determine the slope for the second loading portion of the stress strain curve.

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Experimental

Mechanical properties of the alumina fiber, aluminum matrix, and composites were needed to support the modeling methods. Nextel™ fiber and neat aluminum matrix properties were inputs for the continuous Nextel™ composite strength prediction. The rule-of-mixtures and Hashin's methods used experimental stress strain curves and properties for the individual Nextel™ fabric and Saffil paper reinforced composites.

Fiber Properties

Strength measurements for Nextel™ 610 10,000 denier fiber were obtained using a single fiber tension test. The methodology was similar to that employed by Feih et al. [4] in which each fiber was glued onto a paper tab with Loctite® 454. The paper tab was used to ensure a consistent gage length for each fiber tested. Next, the paper tab was gripped in a TA Instruments Q800 DMA and the sides of the paper tab were cut ensuring only the fiber was carrying load during the test. Then the fibers were loaded in tension at the strain rate specified in ASTM D3552 for tensile testing of metal matrix composites (0.001 s"') [5], Load and strain data were recorded for a total of 90 fiber tests. Stresses were calculated assuming that the fiber radius did not vary from sample to sample. This assumption was based upon fiber diameter measurements made using optical microscopy images of a composite cross section. The strength data obtained in this fashion were fitted to a two-parameter Weibull distribution. The resulting fiber properties are summarized in Table I.

Table 1. Nextel™ 610 Fiber Properties Property Weibull mean strength (MPa), a0

Weibull modulus, m Gauge length (mm), L0

Axial modulus (GPa), E Poisson's ratio, v Fiber radius (um), r

Value 3200 4.5

10.95 370 0.22 5.5

Composite Properties

Four types of materials were tested to obtain mechanical properties including the Nextel™ fabric composite, Saffil paper composite, hybrid (Nextel™ fabric and Saffil paper) composite and the neat matrix alloy. These composites were manufactured by stacking reinforcement layers in a die and infiltrating with liquid aluminum, under pressure. Each composite had an aluminum- 2% copper (A12Cu) matrix. The Nextel™ fabric composite consisted of continuously woven alumina fiber in an eight harness satin weave geometry with a fiber volume fraction of 30%. The Saffil paper composite was composed of randomly oriented short fiber alumina and a 5% volume fraction. The hybrid composite had outer Nextel™ fabric layers with two layers of Saffil paper in the interior. Mechanical properties for the Nextel™ fabric composite, Saffil paper composite, hybrid composite, and neat matrix alloy were measured via uniaxial tension testing. An MTS servo-hydraulic load frame was used to apply strain to rectangular tension samples at a rate of 0.001 s"1 [5]. Alignment of the load frame was within ASTM E 1012 specifications [6]. Strain gages and a digital image correlation system were used to measure strain fields for the sample during loading. Typical stress strain curves for the Nextel™ fabric, Saffil paper, and the hybrid

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reinforced composites are shown in Figure 1. Measured mechanical properties for each composite and the neat matrix alloy are listed in Table II.

Figure 1. Typical Nextel™ fabric, Saffil paper, and hybrid reinforced composites stress strain curves

Table II. Measured composite tensile pro

Composite

Neat A12Cu Saffil Paper Nextel fabric Hybrid

Strength (MPa)

120 140 340 160

Modulus (GPa)

70 85 150 100

perties. Strain to Failure

(%) 12.50 6.36 0.43 0.60

Failure behavior of the hybrid composite was non-catastrophic consisting of an initial failure, reloading, and then ultimate failure, as shown in Figure 2. Point A in Figure 2 represents the initial failure strength which occurred directly before the Nextel™ layer cracked. Once a crack propagated through the Nextel™ layer at point B, the stress dropped by 10% and the composite began to reload. The axial strain fields in Figure 3 were obtained using a digital image correlation system which show the outer Nextel™ layer right before cracking (A) and immediately after a crack has propagated through the layer (B). The cracked hybrid composite carried stress for an additional 0.05% strain until the Saffil layer failed resulting in the ultimate failure of the composite. Fracture profiles of a hybrid composite sample (Figure 4) show multiple cracks which have propagated though the Nextel™ layer but have not entered the Saffil layer.

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Figure 2. Hybrid composite stress strain curve exhibiting multiple failures.

Figure 3. Hybrid composite axial strain field before (A) and immediately after cracking (B).

Figure 4. Hybrid composite fracture profile exhibits cracks through the outer Nextel™ layer.

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Analysis

To predict the hybrid composite's stress strain behavior three methods were utilized. The experimental results of the two base composites were used in a rule-of-mixtures approach to determine the initial loading behavior. Next, Xia and Curtin's [2] continuous fiber reinforced composite strength model was used to predict the cracking strength of the Nextel™ layer through a damage evolution scheme. Finally, Hashin's [3] variational calculus based stress analysis for the hybrid composite containing cracked Nextel™ layers was used to determine the slope for the second loading portion of the stress strain curve.

Continuous Composite Strength Model

Xia and Curtin developed a multiscale damage approach to predict the strength of a continuously reinforced unidirectional metal matrix composite [2], First, stress concentration factors were calculated. A finite element model of the composite at the fiber-scale consisted of a 30-degree wedge section with a central broken fiber and its 12 nearest neighbors. The model was used to simulate stress transfer and calculate stress concentration factors (SCF) in the neighboring fibers. The finite element model defined fibers as linear elastic while the matrix was modeled as elastic-perfectly plastic. The SCFs were then used in a Green's function model which was based on an array of 1024 fibers sectioned into 20um lengths. Each fiber length was given a failure strength according to the fiber Weibull parameters. Load was incrementally applied to the fiber bundle until a fiber failure occurred. Then, the stress carried by the broken fiber was transferred into the composite through frictional pullout and redistribution to neighboring intact fibers based on the SCFs. After the stress was transferred, the model incrementally increases the applied stress until 20% of the fibers have fractured resulting in the composite strength. Ultimate failure was chosen to be at 20% because at this stage the composite was in the cascading failure regime in which additional failures do not appreciably increase the calculated composite strength. A similar approach to Xia and Curtin was used here determine the failure strength of the Nextel™ composite.

Cracked Stiffness Reduction

Hashin developed a variational calculus approach to determine the stiffness reduction of a [0,90], composite with intralaminar cracks [3]. This solution was obtained for a composite with uniformly cracked inner 90° layers. Axial tensile membrane loading was applied to the cracked composite allowing for the stress distributions to be calculated in the 0° and 90° layers. Ultimately, the axial stress recovery and cracked composite stiffness were determined. In this work, the Hashin analysis was applied to estimate the stiffness of the hybrid composite after the Nextel™ layers have failed, but before failure occurs in the Saffil layer.

Results

Application of the rule-of-mixtures, continuous composite strength, and cracked stiffness reduction models to the hybrid composite yielded a prediction for the hybrid composite behavior. Modifications to each method were necessary for application to the hybrid composite. The continuous composite strength model was created for a unidirectional composite; however, it was applied to a woven fabric composite. The stiffness reduction model was developed for cracked inner layers while the hybrid composite had cracked outer layers. These modifications

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did not have significant effects on the prediction which has reasonable agreement with the experimental results.

Continuous Composite Strength Model

Xia and Curtin's continuous composite strength model was applied to the Nextel™ fabric composite. Model inputs included Nextel™ 610 fiber properties, A12Cu matrix properties, and composite geometry. The strength model predicted a Nextel™ fabric strength of 330 MPa while tensile testing measured a 340 MPa strength. Using rule of mixtures to determine the stress at which the Nextel™ fabric composite will cause initial failure of the hybrid composite led to a cracking strength of 160 MPa. This is in exceptional agreement with the experimental data for the hybrid composite.

Cracked Stiffness Reductions

Hashin's variational approach for the cracked composite's stiffness was applied to the hybrid composite. The cracked composite stiffness was predicted with the cracked Nextel™ fabric layer inside of the Saffil paper layers. Inputs for this model included the Nextel™ and Saffil composites moduli, Poison's ratio, and geometry values. The resulting hybrid composite cracked modulus was 30 GPa at a constant crack spacing of 5.19 mm. Visual inspection of the hybrid composite samples had crack spacings between 1.0 and 5.5 mm.

Predicted Hybrid Behavior

In order to predict the stress strain behavior of the hybrid composite, rule of mixtures was used for the portion of the curve up to cracking of the Nextel™ layer. The rule of mixtures approach used the Nextel™ fabric and Saffil paper composite experimental stress strain curves to volume average the portion of load carried in each layer and within the hybrid composite. Xia and Curtin's continuous strength model accurately predicted the hybrid composite initial failure strength of 160 MPa. Finally, Hashin's cracked composite modulus was applied to predict the slope of the cracked composite reloading curve. Using these three approaches, die hybrid composite's stress strain behavior was predicted and shown in Figure 5. Ultimate failure was determined by the hybrid composite's experimental strain to failure value.

Figure 5. Experimental and predicted hybrid composite stress strain behavior.

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Discussion

The approach used here to predict the hybrid stress strain curve was in agreement with the experimental results. Rule of mixtures closely follows the experimental result curve for the initial loading of the composite sample. The continuous Nextel™ fabric model accurately predicted the cracked strength of the hybrid composite. Finally, Hashin's cracked composite modulus matched the initial modulus for the second loading curve. To use these approaches on the hybrid composite a few simplifying assumptions were made. First, the continuous strength model only consists of axial fibers while the Nextel™ fabric consists of both axial and transverse fiber directions. To account for this difference the axial volume fraction was used to model the composite and the transverse fibers were neglected. Even though the transverse fibers were omitted the strength model agreed with the experimental results. There were two major differences between Hashin's approach on a cross ply composite and its application to the hybrid composite. First, the cracked Nextel™ layer was on the outside of the composite laminate samples however it was modeled as the inner layer. Secondly, Hashin's solution was derived for linear elastic composites. The Nextel™ and Saffil composites were not linear elastic and had a yield strength of 20 MPa. To use Hashin's theory a linear elastic modulus was calculated up to the initial failure strength. At this point a damage evolution scheme for the hybrid composite after the initial failure has not been determined. Prediction of the hybrid composite behavior is in excellent agreement with the experimental results up to initial failure. Future investigation of the damage scheme following initial failure will improve the prediction of the second loading curve and ultimate failure.

Future Work

Modeling of the hybrid composite stress strain behavior and failure mechanisms is ongoing. The rule-of-mixtures portion of the prediction will be replaced with a discontinuous Saffil and continuous Nextel™ fabric behavior model. This will allow for the hybrid composite behavior to be calculated using only the fiber and matrix properties. At this point Curtin's stress concentration factors for a similar alumina reinforced aluminum composite were used to predict the failure strength of the Nextel™ composite. A finite element model implementing the current fiber and aluminum properties will be developed to determine the SCFs specific to this problem. Additionally, prediction of the composite damage scheme after initial failure will be improved. Hashin's method will be solved for crack spacings that resemble the cracking patterns of test samples. Interactions between the Nextel™ and Saffil layers will be explored and will result in a prediction for the ultimate failure strength of the hybrid composite. Experimentally, surface replication studies will be conducted during the hybrid composite tension tests to observe the formation and spacing of cracks during the experiment. The methods used here have provided a strong foundation for further exploration into the non-catastrophic failure mechanisms of continuous and discontinuous fiber reinforced hybrid metal matrix composites.

References

1. H. Deve, and C. McCullough, "Continuous-Fiber Reinforced Al Composites: A New Generation." JOM, (1995), 33-37.

2. Z. H. Xia, and W. A. Curtin, "Multiscale modeling of damage and failure in aluminum-matrix composites," Composites Science and Technology, 61 (2001), 2247-2257.

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3. Z. Hashin, "Analysis of Cracked Laminates: A Variational Approach," Mechanics of Materials, 4 (1985), 121-136.

4. S. Feih, a. Thraner, and H. Lilholt, "Tensile strength and fracture surface characterization of sized and unsized glass fibers." Journal of Materials Science, vol. 40, 2005, pp. 1615-1623.

5. ASTM Standard D3552, 2002. "Standard Test Method for Tensile Properties of Fiber Reinforced Metal Matrix." ASTM International, West Conshohocken, PA, 2002, DOI: 10.1520/D3552-96R02. wwvv.astm.org.

6. ASTM Standard E1012, 2005. "Standard Practice for Verification of Test Frame and Specimen Alignment Under Tensile and Compressive Axial Force Application." ASTM International, West Conshohocken, PA, 2005, DOI: 10.1520/E1012-05, www.astm.org.

Acknowledgements

The authors thank the Army Research Laboratory for its financial support and CPS Technologies the metal matrix composite fabrication and technical discussions.

Research was sponsored by the Army Research Laboratory and was accomplished under Cooperative Agreement Number W911NF-06-2-0014. The views and conclusions contained in this document are those of the authors and should not be interpreted as representing official policies, either expressed or implied, of the Army Research Laboratory or the U.S. Government. The U.S. Government is authorized to reproduce and distribute reprints for Government purposes notwithstanding any copyright notation heron.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Numerical Simulation of Pressure Infiltration Process for Making Metal Matrix Composites: Effect of Processing Parameters

Bo Wang1'2 and Krishna M. Pillai1

'Laboratory for Flow and Transport in Porous Media, Department of Mechanical Engineering, University of Wisconsin-Milwaukee, Milwaukee, WI 53211, USA

2School of Materials and Metallurgy, Inner Mongolia University of Science and Technology, Baotou, 014010, China P R

Keywords: pressure infiltration process, metal matrix composites, metal solidification, MMCs

Abstract

A three-dimensional finite difference model is developed for simulating the pressure infiltration process (PIP) for making metal matrix composites (MMCs) where liquid metal is injected under pressure into a mold packed with reinforcing fibers. The infiltration of liquid (pure) metal into the fibrous preform under a constant applied pressure is modeled using the Darcy's law after assuming fully saturated flow behind the liquid-metal front. The concomitant solidification and heat transfer processes are modeled using a suitable energy equation. The simulation allows one to study the effect of process parameters on the PIP infiltration process through evolution of flow-front during filling, solidification pattern, and pressure and temperature distributions. In this study, the effects of the inlet (applied) pressure and the gate size on the infiltration process are studied. The simulation results for such metal infiltration into porous preforms can be used to optimize the PIP and other related manufacturing processes (such as squeeze casting) for making MMCs.

Introduction

Metal matrix composites (MMCs) are widely used in the automobile and aerospace industries. The MMCs can be produced using different methods. The pressure infiltration technique is one of the most important methods to effectively synthesize the MMCs that are reinforced with long continuous fibers. The pressure infiltration process for making MMCs has several advantages: 1) It is relatively inexpensive. 2) Tooling is similar to the tooling for the casting process. 3) It is possible to produce a net or near-shape component from a difficult-to-machine material. 4) It is possible to have a very uniform fiber distribution. 5) It is possible to control the volume fraction and orientation of the fibers. Although there are many advantages in PIP for making MMCs, the difficulties in fabrication, and controlling the process and composites quality have been the major reasons for limiting their commercial applications [1]. The proposed simulation hopes to address these concerns. In the pressure infiltration process (PIP), the molten metal or alloy under the applied gas pressure is injected into a porous preform made of carbon fibers, particles, and whiskers, and is often subjected to higher pressure during cooling and solidification. For this manufacturing process, many mathematical models and numerical simulations have been developed over the past decades to predict and understand the many features of the infiltrated composites [2-16]. The first attempt to model PIP can be attributed to Mortensen [2], where the traditional equations for flow in porous media were employed to develop an analytical model to describe fluid flow, solidification and heat transfer during the infiltration of fibrous preforms by a pure metal.

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Thereafter, the Mortensen research group did several studies on PIP by means of simulation and experimental methods [2-7]. The resin transfer molding, a process similar to PIP, is used for manufacturing polymer matrix composites. In this process, the rate of infiltration and the position of infiltration front are very important for controlling the infiltration process in terms of reducing the formation of dry spots and voids in the fiber preform, and thus improving the quality of the final composite part. Several mold-filling simulations based on the quasi-steady-flow assumption behind a moving resin front in the fibrous preform have been developed to track the flow front [8, 9], The Volume of Fluid (VOF) method has been employed in some studies to model the displacement of the metal/air interface in a fluid environment [10, 11]. The same VOF method as well as the quasi-steady-state assumption behind a sharp front will be employed in the presented simulation as well. When the PIP infiltration is not conducted at a constant elevated temperature, the fiber-metal heat transfer must be taken into account, especially when the fibrous preform and the mold are preheated to a temperature lower than the metal solidification temperature. Several simulations employing this physics have been proposed in the literature. Lacoste et al. [12,13] presented a control volume formulation based on the enthalpy method to simulate the infiltration of aluminum into a preform. Tong and Khan [14] modeled the infiltration by coupling it with solidification and superheat-induced remelting of a pure metal in a porous preform using the finite difference method. Lee and Rohatgi [15] developed a two dimensional model to investigate the solidification process in MMCs when the extended ends of fibers are cooled by a heat sink. Recently, Jung [16] developed a finite element simulation of the squeeze casting process for making MMCs. The slug-flow model with the single-phase Darcy flow behind a clear front was adopted for the simulation and the enthalpy method was employed to predict phase change during the injection of liquid metals in fiber preforms. The objective of this research is to develop a three-dimensional finite difference model to investigate the kinetics of melt infiltration and solidification in PIP under an applied constant pressure. We aim to study the effects of the PIP processing parameters (such as the injection pressure and the inlet-gate size) on infiltration and solidification in this work. A clear understanding of the role of these parameters is necessary to optimize the processing of MMCs using PIP.

Mathematical Model Flow and Solidification Physics

During PIP, the molten metal flows through the fiber preform under applied pressure. The fiber preform, placed inside the PIP mold cavity, is made from reinforcing fibers and is assumed to be a single-scale1 porous medium. Since the molten metal is incompressible, we can express the continuity equation as

du dv ôw „ ,,* — + — + — = 0 (1) ôx dy öz

where u, v, w are the x, y and z components of the volume-averaged Darcy velocity. It has been observed that the liquid flow through porous media can be predicted by the Darcy's law, which agrees with experimental results for low pore-based Reynolds number (defined as Re = uKm Iß where K is the characteristic permeability of a porous medium, u is the Darcy velocity and ft is

1 A porous medium is said to be of the single-scale type when its pores size distribution is unimodal, i.e., the pores are more or less of the same size.

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the viscosity). Because of the low Reynolds-number flows observed in PIP (Re<l), the Darcy's law has also become the most commonly used equation for describing the flow within the PIP mold, and is used as the momentum equation. This equation can be written in a three-dimensional form as

Kx dP K SP K, dP , - . u = - — , v = - — , w = — (2)

/J dx fi dy fi dz

where ju is the viscosity of liquid metal, P is the pore-averaged pressure, and Kx, Ky, Kz are the x, y and z direction (principal) components of the permeability tensor of the fibrous preform2. The Darcy's law applies to the inertia-less, single-phase flow of Newtonian fluids in porous media and is employed here to model the 'plug' flow of liquid metal behind a sharp interface inside the PIP mold cavity. Substituting the Darcy's law, Eqn. (2), into the continuity equation, Eqn. (1), results in an elliptic equation for pressure, Eqn. (3), that can be used to solve the pressure field behind the moving front:

" dx2 ' dy2 ' dz2 W

For tracking the flow-front progression during the mold-filling process, the Volume Of Fluid (VOF) free surface technique is used:

dF udF v dF vi dF „ ... — + + + = 0 (4) dt q> dx q> dy <p dz

Here F represents the volume fraction of fluid contained in a given control volume3. (F = 0 would indicate an empty control volume, F = 1 would indicate a control volume completely-filled by the liquid, and 0 < F < 1 would indicate a partially filled control volume. Note that the front location was given by F= 0.5 .) <p is the porosity of the fibrous preform. The infiltration process can be split into several time steps, and in each time step, the metal flow behind a time-dependent front can be treated as a quasi-steady-state process. It is assumed that the timescale for heat exchange between the fibers and the metal is very short compared to the timescale for infiltration. Hence local thermal equilibrium between fibers and metal is assumed. Furthermore, the viscous dissipation of heat within the metal flowing through small pores of the preform can be neglected. The time-dependent energy equation to predict average temperature in the preform saturated with the molten metal behind the front is given as

r , s „ v ,i5T , - , ST BT ST. . ,S2T d*T d7T, „ Bf, /c\

where p is density, c is the specific heat at constant pressure, T is the average temperature, k is the thermal conductivity of composite, Mim is the latent heat of metal matrix for the solid/liquid phase change, and f, is the solid fraction defined as the ratio of volume of solidified metal to the total volume of metal in the saturated region. The subscripts m, p and c refer to properties of the metal matrix, preform and composite, respectively. Preform permeability is estimated using the equation according to the reference [2]:

tcJjK^IKr ( 6) During solidification, it is assumed that the metal near the reinforced fiber solidifies first and the fiber radius grows as the solidification progresses, thereby causing the effective porosity of the

2 Here we are also assuming that the principal directions of the permeability tensor are aligned with the x, y, z coordinate directions. 3 The whole PIP mold cavity, discretized as shown in Fig. l(Left), is made up of nodes of the finite difference mesh such that control volumes are defined around such nodes.

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preform to decrease. The volume fraction of solid material Vf and the new fiber radius r4 can be calculated using the solid fraction /, as

Kf = Vf+0-vr)f, (7) V = W ^ / K / W where rf is the fiber radius and vf is the volume fraction of fibers in the porous medium. According to [2], for heat flow in the direction perpendicular to the axes of parallel continuous fibers arranged as a square array, the thermal conductivity of the composite medium is given as

* "(kf + km)-{kr-km)e

where kc, *„ and kf are the thermal conductivities of the composite, metal, and fiber, respectively.

Boundary and initial conditions

The boundary conditions are divided into three types based on the region of their application: for the mold walls, for the inlet gates, and for the flow front. At the mold walls, there is no flow in

the normal direction and the condition is expressed as: — = 0. For the virtually insulated mold dn

walls, we can set the adiabatic boundary condition at the mold wall, i.e., — = o. At the inlet gate, on

an applied constant pressure is set at the inlet node: P=PU, and the inlet temperature of melt is set to: T = Tu. At the flow front, the pressure may be set equal to a zero (gage pressure) value or a specified vacuum pressure: P = 0 or P = Pvac.

Problem Description and Numerical Solution

The physical domain of the PIP mold used in our simulation is shown schematically in Fig 1. A three-dimensional cubic mold is used where the size of the mold cavity is 0.2mx0.1mx0.05m. The geometric details of the inlet-gates considered in our parametric study are shown in Fig. 1 (Right). The mold walls are assumed to be adiabatic. The molten metal is driven to infiltrate the porous preform placed in the mold by an externally applied pressure at the mold inlet. The fiber preform considered in the present study is assumed to be homogeneous, which means that its porosity is constant. Furthermore, it is assumed to be made from randomly laid-out fibers; as a result, the principal permeability components (Kx, Ky and Kt) of the preform are equal, i.e.,

Fig. 1 (Left) 3-D finite difference grid used to predict metal infiltration in the PIP preform. (Right) Geometries of the inlet gate considered for the three cases.

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For the present problem, a 22x12x7 (XxY*Z) grid was selected to describe the whole computational domain in a Cartesian coordinate system. An iterative algorithm was developed to solve the flow and solidification equations. At each time step, Eqn. (3) was first solved to estimate the pressure distribution within the wetted region of the preform. Then the velocity distribution was determined by solving the Darcy's law, Eqn. (2). This velocity distribution was then used to solve for the temperature using Eqn. (5). The velocity was also used to advance the flow front using Eqn. (4). The solution process was repeated in the next time-step. The discretized form of the conservation equations were solved using our Fortran code MIMPS° (Metal IMpregnation Process Simulation).

Results and discussion

In this study, the molten aluminum is pressurized by a constant inlet pressure and is made to infiltrate into the PIP mold cavity packed with the fiber preform. The physical properties of the pure melt and reinforcement fibers are presented in Table 1. The initial temperatures of the melt, the mold, and the preform are 1000K, 730K, and 730K, respectively. Table 2 lists the process parameters for this simulation. The effect of inlet-gate size on infiltration is studied by changing the size and location of the gate as shown in Fig. l(Right).

Table 1. Material properties

Al(l) Al(s) Carbon Fiber

ptkg/m1) 2400 2700 3300

Cf(JIKg-K)

1083 1296 1213

H(l'a-s)

1.3xl0'3

k(Wlm-K)

93

8

T,(.K)

933

Ah^kJIkg)

398

Table 2. Process parameters for the present simulation Pnrantfftorv vnhif Parameters Inlet Pressure, Plrl

Inlet temperature of the melt, TM

Preheat temperature of preform, T

Preheat temperature of mold, r„

value 3 atm

1000 K

730 K

730 K

We firstly discuss the pressure infiltration process for a situation with the gate dimension the same as the width of the mold wall. The time variable for the infiltration process is made dimensionless as t' =tltm where tm is the time needed to fill the PIP mold. Fig. 2 shows the evolution of the metal flow-front with time for this situation. The front moves forward uniformly and reaches the other side of mold at the end of infiltration. The applied inlet pressure will be maintained constant in PIP so that the preform can be fully saturated and the gas pockets can be completely dissolved at the end of the infiltration process. Fig. 3 presents the distribution of pressure and velocity at the central x-y section of the mold at various times. From figures 2 and 3, it can be seen that the flow front at each time tends to be planar since the fiber preform is isotropic and the gate dimension is equal to the width of the mold. The pressure distribution is uniform along the direction perpendicular to the infiltration direction, and the infiltration velocity decreases rapidly with time. The temperature distribution for this case can be seen in Fig. 4. Due to the adiabatic boundary condition imposed at all side mold-walls, the temperature gradient is concentrated near the flow front. Similar to the pressure and velocity distributions, the

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temperature distribution is also uniform along the direction perpendicular to the infiltration direction and heat transfer mainly happens along the flow direction.

(a) (b) (c) Fig. 2. Evolution of metal metal-front inside the considered PIP preform at different dimensionless times: (a) 0.015, (b) 0.09, and (c) 0.92.

(a) (b) (c) Fig. 3. The pressure distribution and velocity field in the central plane at the different dimensionless times: (a) 0.015, (b) 0.09, and (c) 0.92.

(a) (b) (c) Fig. 4. The temperature distribution in the central plane at different dimensionless times: 0.015, (b) 0.09, and (c) 0.92.

(a)

Inlet applied pressure, atm

Fig. 5. Effect of the inlet (applied) pressure on the infiltration time.

Fig. 5 shows the effect of the inlet (applied) pressure on the total infiltration time. The total infiltration time is defined as the time interval from the beginning of infiltration to the instant the infiltration front reaches the right wall of the mold. The simulation results show that the infiltration time decreases monotonically with an increase in the inlet pressure. But as the inlet pressure reaches a higher value, the decrease in the infiltration time becomes progressively smaller.

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Fig. 6 shows the pattern of flow front evolution with time for the remaining two inlet-gates. Fig. 7 presents the temperature distribution at dimensionless time = 0.96 for the remaining two inlet-gate openings. It can be seen from figures 2, 4, 6 and 7 that the shapes of flow fronts as well as the distributions of flow variables such as temperature depend strongly on the inlet gate opening. However when the front meets the mold walls (top and bottom side-walls), the confinement of the flow within parallel walls makes the filling behavior unidirectional and flow front tends to be planar at higher infiltrating times. (When the dimensionless infiltration time is greater than 0.54 for Fig.6-a and greater than 0.79 for Fig.6-b, the shape of flow front has become almost straight.) Also it is obvious that the temperature distributions replicate the behavior of the flow fronts. It is also noted that the extent of the zone where the temperature gradient is concentrated increases. This kind of changes in temperature distributions is likely to have an important effect on the development of the grain-level and dendrite-level microstructures in the final MMC parts made using PIP.

Fig. 6. The evolution of flow front at different infiltration times for the remaining two kinds of gate-openings: (a) gate located at the center of the left mold-wall, and (b) gate located on the lower side of the left mold-wall. (The x and y axes values are in meters while the time associated with various contours is the dimensionless time.)

Fig. 7. Temperature distribution at dimensionless time = 0.96 for the remaining two kinds of gate-openings: (a) gate located at the center of the left mold-wall, and (b) gate located on the lower side of the left mold-wall.

Summary and Conclusions

A numerical simulation of the infiltration of a pure metal Aluminum in a three-dimensional PIP mold using the homogeneous and isotropic fibrous preform is presented in this paper. The infiltration and solidification process are predicted during the mold-filling simulation. The results

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show that the flow-front shape and the distribution of variables such as temperature depend strongly on the inlet-gate location and size. The mold fill-time is shown to decrease monotonically with the inlet pressure. These results will be helpful for understanding the flow and transport phenomena occurring in the practical PIP, and thus correlate with the experimental observations. It can also be used for optimizing the PIP molds used for making MMCs.

Acknowledgments

The authors would like to acknowledge the financial support received from the Research Growth Initiative (RG1) program of the Graduate School at University of Wisconsin-Milwaukee and the Catalyst Grant offered by the Bradley Foundation. We also wish to thank the help and comments received from Professor Rohatgi's group for this research. The first author also gratefully acknowledges the financial support received from the Open Project of Shanghai Key Laboratory of Modern Metallurgy and Material Processing (SELF-2010-04) and the Project for Innovation Fund of Inner Mongolia University of Science and Technology (No.2009NC006).

References

1. ChawlaN, Chawla KK. Metal Matrix Composites (New York, NY: Springer-Verlag, 2005), 351-401. 2. Mortrensen A et al., "Infiltration of fibrous preforms by a pure metal Part 1 theory," Metall Trans A, 20A(11)( 1989), 2535-47. 3] Mortensen A, Michaud V, "Infiltration of fiber preforms by a binary alloy: Part I. Theory," Metall Trans A, 21A(7)(1990), 2059-72. 4. Andrews RM, Mortensen A, "Lorentz force infiltration of fibrous preforms," Metall Trans A, 22A(12)(1991), 2903-15. 5. Jarry P et al., "Infiltration of fiebr preforms by an alloy: Part 3 Die-casting experiments," Metall Trans /!, 23A(8)(1992), 2281-9. 6. Michaud VJ, Mortensen A, "Infiltration of fiber preforms by a binary alloy: Part 2 Further Theory and experiments," Metall Trans A, 23A (8)(1992), 2263-80. 7. Michaud VJ, Sommer JL, and Mortensen A, "Infiltration of fibrous preforms by a pure metal: Part V. Influence of preform compressibility," Metall Mater Trans A, 30(2)(1999), 471-82. 8. Young WB, "3-dimensional nonisothermal mold filling simulations in resin transfer molding," Polym Compos, 15(2)(1994), 118-27. 9. Shojaei A, Ghaffarian SR, and Karimian SMH, "Numerical simulation of three-dimensional mold filling process in resin transfer molding using quasi-steady state and partial saturation formulations," Compos Sei Technol, 62(6)(2002), 861-79. 10. Voller VR, Peng S, "An algorithm for analysis of polymer filling of molds," Polym Eng Sei, 35(22)(1995), 1758-65. 11. Lin M, Hahn HT, Huh H, "A finite element simulation of resin transfer molding based on partial nodal saturation and implicit time integration," Composites Part A, 29(5-6)(1998), 541-50. 12. Lacoste E et al., "Numerical-simulation of the injection-molding of thin parts by liquid-metal infiltration of fibrous preform," Mater Sei Eng A, 135(1991), 45-9. 13. Lacoste E et al., "Numerical-simulation of the infiltration of fibrous preforms by a pure metal," Metall Trans A, 24(12)(1993), 2667-78. 14. Tong X, Khan JA, "Infiltration and solidification remelting of a pure metal in a two-dimensional porous preform,'' J Heat Transf-Trans ASME, 118(1)(1996), 173-80. 15. Lee E-K, Amano RS, Rohatgi PK, "Metal matrix composite solidification in the presence of cooled fibers: numerical simulation and experimental observation," Heat and Mass Transfer, 43(8)(2006), 741-8. 16. Jung CK, Jang JH, Han KS, "Numerical simulation of infiltration and solidification processes for squeeze cast al composites with parametric study," Metall Mater Trans A, 39A( 11 )(2008), 2736-48.

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TIMIS2012 141s t Annual Meeting & Exhibition

Titanium: Advances in Processing, Characterization

and Properties

The proceedings contained in this section have not been edited or reviewed by the conference program organizers. The opinions and statements expressed in the proceedings are those of the authors only and are not necessarily those of the editors or TMS staff. No confirmations or endorsements are intended or implied.

Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Microstructural evolution and mechanical properties of ß-titanium Ti-10V-2Fe-3Al during incremental forming

S. Winter1, S. Fritsch1, M.F.-X. Wagner1

Chemnitz University of Technology; Erfenschlager Str. 73, Chemnitz, 09117, Germany

Keywords: Ti-10V-2Fe-3Al, beta titanium, incremental forming

Abstract

Forming of hollow shafts of high-strength beta titanium alloys is technologically demanding, as well as expensive. One approach to reduce costs is incremental forming by spin extrusion. The advantage of this process is the high utilization of material (80 %) in comparison to deep-hole drilling (40 %). In this study, we investigate the microstructural evolution and the mechanical properties of a Ti-10V-2Fe-3Al hollow shaft formed by spin extrusion, which results in varying stress-strain behavior across the wall thickness. Grain refinement from >1 urn (edge) to 100 um (center) is observed. Subsequent heat treatments can be used to achieve homogeneous properties throughout the shaft's cross section. Even higher strengths are associated with the precipitation of primary a-phase. Our results for the high-strength beta titanium alloy Ti-10V-2Fe-3Al illustrate the potential of spin extrusion, combined with suitable heat treatments, to produce hollow shafts with improved properties at lower costs.

1. Introduction

Excellent mechanical and corrosion properties are the main reasons for choosing titanium materials, primarily in the aerospace industry [1, 4, 9]. The properties of different alloys are strongly affected by the fractions of hexagonal close-packed (hep) a-phase and body-centered cubic (bec) ß-phase. Compared to the a-phase, the ß-phase is characterized by a lower strength and a higher formability. A combination of good ductility and fatigue resistance makes ß-titanium alloys interesting for many current and potential future applications. The mechanical behavior of titanium alloys is primarily determined by microstructural parameters like grain size (ß-phase), morphology and volume fraction of primary / secondary a-phase precipitates and this allows changing and optimizing their mechanical properties across a wide range [1, 5, 7, 10]. However, ß-titanium alloys are quite expensive compared to other materials, and in the past this issue has limited their actual use considerably. Titanium hollow shafts are usually produced by cold extrusion. This technique produces work-pieces within fine tolerances and with excellent surface topographies. However, the process is limited to thin-walled tubes since it is restricted to certain length to diameter ratios [2]. The alternative process of deep-hole drilling has a material utilization of only 40 %, which makes it very inefficient for high-grade materials [6]. Recently, another alternative, based on incremental forming steps, was developed at the Fraunhofer Institute for Machine Tools and Forming Technology (Institut fur Werkzeugmaschinen und Umformtechnik, IWU) in Chemnitz. The so-called spin extrusion can be described as a partially rotating forming technology to produce axially symmetric hollow shafts directly from solid semi-finished bars. The hollow shafts are created using a mandrel that shapes the inner contour, and spinning rollers that rotate along the outer surface of the semi-finished bar, as illustrated in Figure 1. Both components are moving simultaneously in axial direction, and therefore the bar's material yields axially against the feed direction, forming a cup wall. The work-piece is chucked in the spindle and rotates around its

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longitudinal axis; rollers, punch and spindle rotate synchronously [6]. The major advantage of spin extrusion over deep-hole drilling is the significantly higher utilization of material (nearly 80 %) [3, 6], which potentially allows to produce ß-titanium hollow shafts at lower costs. At present, there is a lack of experimental data regarding the properties of the ß-titanium alloy Ti-10V-2Fe-3Al formed by spin extrusion. One goal of this study is to demonstrate how the microstructural evolution during incremental forming influences the mechanical properties, in particular as a function of the location in the cross section of the hollow shaft.

Figure 1. Schematic diagram of the spin extrusion process and the relevant parts and components involved in the incremental forming process.

2. Materials and Experiments

Rods of Ti-10V-2Fe-3Al (Ti-10-2-3) were obtained from Otto Fuchs KG in a solution-annealed condition. In this heat treatment condition, the material exhibits a yield strength (YS) of 620 MPa, an ultimate tensile strength (UTS) of 815 MPa, a uniform elongation of 32 % and an elongation at fracture of 35 %. The material is characterized by a large ß-grain size of approx. 300 |im, as shown in the optical micrograph in Figure 2.

Figure 2. The microstructure of Ti-10V2Fe-3Al in the solution annealed condition exhibits large grains.

Forming of the hollow shafts by spin extrusion was performed by IWU in Chemnitz. Previous work demonstrated that forming of Ti-10-2-3 can only be performed in a rather small temperature window between 500 and 600 °C. Lower forming temperatures generally lead to defects in the work-piece, such as cracks; at higher temperatures, the material strength decreases considerably, and the forming process also becomes prone to failures [6], Therefore, the shafts that were subsequently analyzed in this study were formed at IWU in the appropriate temperature range given above.

834

For hardness measurements (Vickers hardness: HV30) across entire cross-sections of the shaft, it was cut into eight rings, and about 250 indentations into the surface of each ring were carried out. Furthermore, small ("mini") tensile specimens were taken out of the ring walls, Fig. 3. Depending on the wall thickness, 4 to 6 specimens could be taken out of individual rings. The tensile axes of the flat specimens were oriented parallel to the central axis of the shaft. Therefore, a systematic characterization of the stress-strain behavior as a function of the distance to the outer surface of the shaft could be performed. For tensile testing, we used a UPM 1475 Zwick/Roell machine with an extensometer clipped onto the gage length (10 mm). The uniaxial tensile tests were performed at an initial strain rate of 10"3 1/s.

Figure 3. Sampling of tensile specimens out of a ring. The tensile axes are oriented parallel to the central axis of the hollow shaft.

Some of the shaft's material was subjected to an additional heat treatment that consisted of three steps: In the first step, the material was annealed at 850 °C for 4 min in order to achieve recrystallization of the microstructure. The second and third steps were carried out to precipitate the primary and secondary a-phases. For precipitation of the primary a-phase, the material was heated to 725 °C for 1 hour, and then annealed at 540 °C for 8 hours. This leads to the formation of the high-strength, secondary a-phase. On the heat treated materials, hardness measurements and tensile tests similar to the ones described above were performed, and the results were compared to those of the material after spin extrusion in order to document the changes of mechanical properties due to the heat treatments. The microstructures of all materials were analyzed by complementary optical microscopy and scanning electron microscopy (SEM).

3. Results and Discussion

Typical results of the hardness measurements on the as-spin-extruded material are shown in Figure 4, where the distribution of hardness values is shown across one ring. An inhomogeneous distribution of hardness values can be observed, with differences of more than 60 HV30 as one moves from an inner or outer region (with hardnesses of about 350 HV30) to the center line of the ring (with about 300 HV30).

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Figure 4. Distribution of hardness values across the shaft's cross-section after spin extrusion. Regions near the inner and outer surfaces of the ring exhibit hardness values between 330 and

360 HV30; the central region of the ring shows hardnesses between 300 and 315 HV30.

The complex deformation during the spin extrusion processing also results in a variations of the stress-strain behavior across the wall thickness, Fig. 5. Because of the high degree of deformation, there is no significantly strain hardening of the samples, and the ultimate tensile strength is only slightly higher than the yield strength in all cases. The engineering stress-strain curves qualitatively agree well with the hardness measurements: specimens taken from the center of the ring exhibit the lowest strengths (YS of 940 MPa), whereas the outer regions show yield strengths between 1000 and 1100 MPa. We note that this represents a considerable increase of yield strength (300 to 400 MPa) compared to the solution annealed material. This change in mechanical properties results from pre-heating the rods to 650 °C (for 1 hour) before the forming process itself, which leads to precipitation of the strength- and hardness-enhancing primary a-phase [1,5].

Figure 5. Representative stress-strain behavior after spin extrusion at different locations across the ring's wall.

The differences in terms of both hardness distributions and uniaxial tensile behavior across the cross-section of the hollow shaft can be related different deformation regions and cooling rates in the forming process of the shaft: The material near the inner and outer surfaces is highly deformed, whereas the central regions are clearly subjected to smaller amounts of plastic deformation. This observation is also in line with the different microstructural features in the different regions of the material, Fig. 6. The grain size in the highly deformed regions is about 1 to 5 urn, and it increases strongly towards the less deformed regions in the center part of the

836

cross-section. No attempt was made to determine the grain size distribution quantitatively, but even from our qualitative observations it is obvious that there are quite huge, elongated ß-grains with an approx. grain size between 100 and 300 um (i.e., somewhat similar to the grains observed in the solution treated initial state, see Fig. 2). Clearly, the grain refinement in the inner and outer regions of the ring is one of the main reasons for higher strength and hardness in comparison to the center.

Figure 6. Microstructure of the hollow shaft after spin extrusion: (a) coarse, elongated grains are clearly visible in the center (grain sizes 100 - 300 |im) and fine grains (1-5 urn) are found in the

inner and outer regions, as shown with increased magnification in (b).

One additional reason for the differences in mechanical properties is an increased amount of the primary a-phase in the inner and outer regions. This is evident from the SEM micrographs shown in Figure 7. The globular, hep a-phase has a higher strength then the ß-phase matrix, and therefore a higher amount of a-phase is directly related to an increase of strength [1, 5, 10].

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Figure 7. Volume fractions of strength increasing globular a-phase in different areas of the shaft: (a) large volume fraction in the outer area; (b) lower volume fraction in the central region.

While an in-depth analysis of the relation between processing parameters, microstructural and mechanical inhomogeneity is scientifically interesting (and further, quantitative results will be published elsewhere), inhomogeneous mechanical properties are detrimental for most potential future applications. Therefore, we now focus on the effect of the subsequent heat treatment. We highlight again that this heat treatment was introduced with the aim of homogenizing the microstructure and mechanical properties of the titanium hollow shafts. The hardness data shown in Figure 8 demonstrates that hardness values are much more homogeneously distributed throughout the cross-section compared to the as spin extruded material. The average hardness is 315Ü5HV30.

Figure 8. Homogeneous hardness distribution after the heat treatment, with an average hardness ofabout315±15HV30.

Similarly, after the heat treatment, the stress-strain behavior becomes quite homogeneous, Figure 9. All stress-strain curves correspond to the same level of strength. At 980 MPa throughout the whole cross section, the yield strength is slightly lower than the strength of the untreated rings. Despite the heat treatment, there still is no significant strain hardening, and therefore the ultimate tensile strength is only slightly higher than the yield point. In contrast to the untreated material, however, the stress-strain curves to exhibit rather extended strength plateaus. Relating our mechanical results to me microstructural evolution during our heat treatment, we note that after recrystallization, both the primary and secondary a-phase are precipitated. The secondary a-phase leads to higher strength than the primary phase [1, 8, 10],

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and hence the strength of the heat-treated material is on a level similar to that of the highly deformed hollow shafts prior to our heat treatment.

Figure 9. Representative, homogeneous stress-strain behavior after the subsequent heat treatment.

Figure 10. Homogeneous microstructure after the subsequent heat treatment with a grain size of about 45 um (optical micrograph).

Finally, as is to be expected from the mechanical properties, the microstructure becomes homogenized by the heat treatment. Figure 10 shows that the grain size increases (from 5 |im in the inner and outer regions after spin extrusion, see also Figure 6) to 45 urn throughout the entire cross section.

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4. Summary and Conclusions

This paper presents first results of an ongoing investigation of the possibility to produce homogeneous titanium hollow shafts by means of the spin extrusion technique. Hardness measurements and tensile tests with small specimens were conducted. After the incremental forming process, considerable variations in terms of mechanical properties and grain sizes can be observed in the shaft's cross section. The microstructure shows grain refinement from 1 - 5 um (surface regions) to about 100 - 300 um (center). We have proposed a subsequent heat treatment consisting of three steps that was successful in homogenizing the microstructure and mechanical properties (analyzed in terms of hardness distributions and tensile stress-strain curves) of the shaft. As yield strengths and ultimate tensile strengths are significantly increased in comparison to the solution annealed material prior to forming, spin extrusion seems to be an efficient method to produce ß-titanium hollow shafts with improved properties at reduced costs.

Acknowledgements

The authors gratefully acknowledge funding by the German Research Foundation (Deutsche Forschungsgemeinschaft DFG) in the framework of PAK 292.

References

1. M. Peters, C. Leyens, Titan und Titanlegierungen (WILEY-VCH, 2000), 49-68.

2. R. Neugebauer et al., Innovative Cold Forming Processes for Increasing Efficiency of Energy Consumption and Resources Process Chains for Manufacturing of Power Train Components (Düsseldorf: Tagungsband 25. Jahrestreffen der Kaltmassivumformer, 2010).

3. R. Neugebauer et al., "Incremental Forming of Hollow Shapes", Steel research int., 76 (2/3) (2005)

4. Matthew J. Donachie, Titanium - a technical guide (ASM International, 2000).

5. J. Kiese, L. Wagner, "Fatigue behaviour in Ti-10V-2Fe-3Al: Microstructural effects on crack nucleation and microcrack growth", Fatigue '96, Pergamon Press, 1996.

6. R. Neugebauer et al., "Manufacture of a b-Titanium Hollow Shaft by Incremental Forming", Production Engineering, 5 (2011), 227-232.

7. A. Drechsler, Einfluss des Festwalzens auf das Dauerschwingverhalten der metastabilen Beta-Titanlegierung Ti-10V-2Fe-3Al (Fortschritt-Berichte VDI Reihe 5, Nr.633).

8. S. K. Jha, R. Chandran, "Effect of secondary alpha in beta phase on the fatigue of beta titanium alloy Ti-10V-2Fe-3Al", Fatigue 2002, 1815-1822.

9. P. Allen, "Titanium alloy development", Advanced Materials & Processes, 150 (4) (1996), 35.

10. B. Wang et al., "Mircrostructural evolution during aging of Ti-10V-2Fe-3Al titanium alloy", Journal of University of Science and Technology Beijing, 14 (4) (2007), 335.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Study on Hot Deformation Behavior of TC4 Titanium Alloy

Yanling Lu , Sihai Jiao , Xinglai Zhou , Anping Dong

(1. Department of Nuclear Materials Science and Engineering, Shanghai Institute of Applied Physics,

Chinese Academy of Sciences, Shanghai, 201800;

2. Metallurgical Process Department, R&D Center, Baoshan Iron & Steel Co., LTD., Shanghai, 201900;

3. School of Materials Science and Engineering, Shanghai Jiao Tong University, Shanghai, 200240)

Abstract: Hot compression deformation behavior of TC4 titanium alloy was studied on Thermecmaster-Z

simulator in the temperature range SSO"^! 150°Cand strain rate range 0.5 s"1—30s"1. The results show that

deformation temperature and strain rate both have significant influence on the flow stress. The flow stress

decreases with the increase of deformation temperature, while increases with the increase of strain rate. The

deformation mechanism of TC4 alloy exhibit dynamic recovery feature in high temperature. However, in

lower temperature, dynamic recrystallization and grain boundary slip behavior may take place.

Deformation activation energy values are 863.46 kJ/mol in (a + ß) phase region, and 209.80 kJ/mol in ß

region respectively. Constitutive equations of TC4 alloy described by Zener-Hollomon parameter were

formulated. Therefore, a scientific basis is provided for the reasonable choice of thermal parameters of TC4

titanium alloy.

Key words: TC4 titanium alloy; Hot compression deformation; Flow stress; Constitutive equation

1. Introduction

TC4 (Ti-6A1-4V) alloy is an important two phase (a+ß) titanium alloy including 6% aluminum and

4% vanadium. TC4 alloy has extensively used in aerospace industry due to its high relatively strength and

good resistance against heat and corrosion [1-3].

Recently, FE (Finite Element) simulation has become one of the most important and effective methods

to analyze material deformation processing. As we know, the value of flow stress has significant influence

on the simulation accuracy. Flow stress is one of the basic parameter to material plastic deformation

depending on the deformation temperature, strain rate and chemical composition. Hot deformation

processing is generally used for TC4 alloy because of its high deformation stress [4,5].

In this paper, effect of the deformation temperature, the strain rate and the strain on the flow stress is

analyzed to represent the mechanical behavior of TC4 alloy, and the activation energy for deformation of

compressed TC4 alloy is calculated. Besides, the deformation mechanisms are clarified through the

activation energy in (a + ß) phase region and the single ß phase region. Finally, Constitutive equation of

TC4 alloy described by Zener-Hollomon parameter is established so as to optimize the processing

parameters.

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2. Experimental procedure

The experimental material is produced at the Special Steel Branch, Bao Steel of China. The chemical

composition of the experimental TC4 titanium alloy is given in Table. 1. The as-received microstructure of

the TC4 alloy consisted of a phase and ß phase is illustrated in Fig.l. It can be seen that the microstructure

is the fine equiaxed-shaped a grains having the grain size lOum. According to calculation, the phase

transformation temperature of TC4 alloy is about 975'C.

Figure. 1 Micrograph of the as-received TC4 titanium alloy

Cylindrical compression specimens were machined from the hot-rolled bar with 50mm in diameter to

have dimensions of 8mm in diameter and 12mm in height. Hot compression deformation behavior of TC4

titanium alloy was studied on Thermecmaster-Z simulator. The test temperature ranged from 850"C to

1150'C and strain rates were kept at 0.5 s"1, 5 s"1, and 30s"1. The specimens were deformed to about 50%,

then the specimens were cooled to room temperature at the 10*C/s cooling rate.

3. Results and discussion

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(a) (b)

Figure.2 Effect of deformation temperature and strain rate on flow stress in TC4 titanium alloy (a) Strain rate= 30s"1; (b) T=1000'C

3.1 Effect of Deformation Parameters on Flow Stress

To investigate the flow behavior in (a + ß) phase and the single phase (ß phase) regions, the typical

stress-strain curves at a strain rate 30s"1 and deformation temperature 1000t are given in Fig.2. Under all

experiments, the flow behavior of TC4 alloy exhibits an initial increase in the flow stress with increasing

strain until a peak stress value beyond which there is a decrease. From Fig.2, it can be seen that

deformation temperature and strain rate both have significant influence on the flow stress. The flow stress

decreases with the increase of deformation temperature, while increases with the increase of strain rate. At

the strain rate 30s"1, when the deformation temperature increases from 850'C to 1150'C, the peak stress

decreases from 339.85 MPa to 72.38MPa. Similarly, in the case of hot deformation temperature 1000'C,

when the strain rate increases from 0.5 s'1 to 30 s"1, the peak stress increases from 45.01 MPa to 112.62

MPa correspondingly. That is, the flow stress of hot compression TC4 alloy is significantly sensitive to the

deformation temperature and the strain rate. With the increasing strain rate, dynamic recovery and

recrystallization become more and more difficult so as to worsen the alloy plastic properties and flow

stress increases conversely.

Fig.2 illustrates that at the range of experimental deformation temperature and strain rate, the shape of

strain-stress curves is different, however, there is peak of flow stress occurred. Besides, the peak moves

ahead with the decreases of deformation temperature. From Fig.2, it can be seen that the softening effect

in the two phase (a+ß) region is significantly different from that in the ß phase region. In the two phase

(a+ß) region, the flow stress increases quickly with the increasing of strain and reaches a peak value.

Then the flow stress sharply decreases, that is, there is an apparent stress peak resulting from dynamic

recrystallisation. However, in the single ß phase region, the strain-stress curve becomes gradually steady

resulting from dynamic recovery. Wanjiara et al [6] also observed that the differences of flow softening

behavior in the ß phase region with that in the a + ß region occurred in compression of near type titanium

alloy IM1813. Above a deformation temperature of 1000'C, TC4 alloy belongs to the ß single-phase with

843

bec structure, so inter-slip of dislocation may occur easily.

3.2 Constitutive Equation

High temperature deformation of the metal is a hot activation process. Strain rate and deformation

temperature have the important effects on the flow stress, and the relationship can be depicted by Arrhenius

equation [7-9]:

© For the strain value smaller than 0.8 ( aa <0.8 ) :

s = Axa" ( 1 )

© For the strain value higher than 1.2 ( acr>\.2):

È = A2 exp(/?cr) (2)

® For the whole strain range:

è = 4sinh(aer)]" exp[-Q/(RT)] (3)

where éis the strain rate, A\, Ai, A, a, and ß are constant, cris the stress peak value, n is the strain

constant» Q is the deformation activation energy, R is the gas constant, and Tis the absolute deformation

temperature. The relationship of a. ß% n can be written as followed:

a = ßln

Equations (4)and (5) can be calculated by equations ( 1 ) and (2):

lnê = In A[ + »In er (4)

In ê = In A2 + ßa- (5)

By using regressive method, two curves of In er - In ê and a - In s are obtained, as shown in Fig.3.

It is seen that the two curves of In a — In ê and a - In s of TC4 alloy is in liner relationship when the

deformation temperature varied from 850°C to 1150°C n and ß are the reciprocal of linear slope of Fig. 3(a)

and (b) according to equations (4) and (5) respectively. It is can be seen from the two curves of

In er - In È and er - In s , the slope has a big difference compared high temperature with low

temperature when the phase transformation point is dividing line. The average value are as followed: n i„w

iemp«ature=9.04736, n high cemp«roture=5.49504, ßi„wtemperauire=0.04997, ßhigh tempcraturc=0.09139. According to

the equation or - ßl n, the values of a in low temperature and higher temperature are 0.0055 and 0.0166,

respectively.

Then following equation (6) can be obtained from taking logarithm of the equation (3):

lnff = ln/* + «ln[sinh(aer)]-g/(/?7') (6)

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The relationship of In è - ln[sinh(orer)] is shown in Fig. 4, Ins - ln[sinh(o;<7)] is almost in linear

relationship. The average reciprocal of linear slope after regressive are as below:

* low temperature O . j U Z j Z , A high temperature ^- l ^ o O l

(3) (b)

Figure.3 Relationship of peak stress and strain rate for TC4 titanium alloy

Figure.5 Relationship of flow stress and Figure.4 Relationship of flow stress and d e f o r m a t i o n temperature for TC4 titanium alloy strain rate for TC4 titanium alloy

In the equation (6), C, = (In s - In A) I n , C1=Q/(l OOOnR), then the equation (7) is:

ln[sinh(ao-)] = C 1 + C 2 * 1 0 0 0 / r (7)

The relationship of ln[sinh(or<T)] — 1000 / T is shown inFig.5. The Arithmetic average slopes of the

curve are 15.97214 in the high temperature and 6.11203 in the low temperature respectively. It is concluded

that TC4 has the different deformation activate energy when the experiment temperature changed, it is

different with previous research results[10]„

845

Equation (8) is the differential curve of the equation (3):

Q = R[ ^ l ain[smh(g<7)] (

dln[smh(aa)]T d(\IT)

Using the slope value calculated by Fig. 4 and Fig.5 in equation (8), then Q value in the high

temperature and low temperature are as below:

Q=863.46 kJ/mol ( low temperature) , Q=209.80kJ/mol (high temperature)

TC4 alloy has different deformation activation energy in different temperature range, it is illustrated

that the deformation mechanism of TC4 is not the same when the temperature changed. The self-diffusion

activation energy of the o—Ti is 170 kJ/mol. this value for the ß—Ti is 153 kJ/mol[ll] The deformation

activation energy of the ß phase region calculated by the author is close to 153 kJ/mol[ll], it means that the

deformation of TC4 in ß phase region are mainly controlled by self- diffusion. However, the deformation

activation energy in two phase region is much larger than the self-diffusion activation energy of Ti alloy,

which means that the deformation process is controlled by the other mechanism except for self-diffusion. It

is believed that grain boundary sliding is the main deformation mechanism of (a-r-ß)phases region[12, 13].

Based on the previous research, it is true that the hot deformation activation energy of the metal is close to

the self-diffusion activation energy in the case of the dynamic recovery. However, the hot deformation

activation energy is much higher than self-diffusion activation energy in the case of the dynamic

recrystallization. In this paper, the activation energy of the TC4 alloy in the lower temperature range is

863.46 kJ/mol, the results showed that the dynamic recrystallization occurred during the process of

deformation.

3.3 Relationship of Z Parameter and Flow Stress

In 1944, Zener and Hollomon proposed the relationship of strain rate and temperature and investigated

it experimentally.

Zener-Hollomon parameter can be expressed as follows [14]:

Z = s exp(0 IRT) = ̂ [sinh(ao-)]" ( 9 )

Where Z is Zener-Hollomon parameter, è is the strain rate, Q is the hot deformation activation

energy, R is the gas constant, and T is the absolute temperature.

Then equation (10) can be obtained from taking logarithm of the equation (9):

lnZ = l n ^ + «ln[sinh(a<r)] (10)

Z can be obtained by substitution of the Q value in equation (9), the relationship of

In Z - ln[sinh(orcT)] is shown in Fig. 6. It can be found that In Z - ln[sinh(oro')] is in good agreement

with linear relationship, thus it can be further proved that equation (3) is correct in depicting the

relationship of flow stress, strain rate and temperature of hot press deformation of TC4 alloy. So, some

parameters can be obtained in different temperature range according to Fig.6:

In the lower temperature case: n = 6.24149, InA = 88.27142, A =2.167x 1038.

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In the higher temperature case: «=3.84329, ]nA = 19.31316, A =2.441 * 10 .

Figure.6 Relationship of flow stress and Z parameter for TC4 titanium alloy

Equation (11) can be obtained based on equation (8) and the definition of hyperbolic sine function:

c7 = lln((4)"" + [(4)2'"+l]"2j CD a { A A J

Accordingly, constitutive equations of flow stress of TC4 alloy can be expressed by Z parameter:

in (a+ß) two phases area: a=0.0055,/f =2.167xl03S, « = 6.24149

a- = — J — l n ( ( f L _ ^ " « 4 < « + [( Z fw + yi'2 0.0055 l 2 . 1 6 7 x l 0 3 8 2.167xl038

in ßphase area: a=0.0166,^ =2.441xlOs, «=3.84329

a = — ^ — l n l ( 7L-—f<^^ + [( 1 f>w + !]■« 0.0166 1 2.441 x 10s 2.441 x 108

The error of average square root is less than 6% between the calculated flow stress " and the tested

4. Conclusions

( 1 ) Deformation temperature and strain rate have significant effects on the flow stress of TC4 titanium

alloy. The flow stress decreases with the deformation temperature increases and increases with the

increases of strain rate.

(2) The strain-stress curve of TC4 has an obvious peak value when the deformation temperature range

from 750 "C to 950 *C , deformation mechanism is grain boundary slip and dynamic

recrystallization. However, when the temperature is higher than 1000'C, the strain-stress curve

becomes steady, the deformation mechanism of TC4 alloy exhibit dynamic recovery feature.

(3) Deformation activation energy values of TC4 titanium alloy are 863.46 kj/mol in (a + ß) phase

region, and 209.80 kJ/mol in ß region respectively.

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(4) Constitutive equations of TC4 alloy is as followed:

In a + ß phase region:

a = —!—lJ( 2 - T T )""4149 +1( Z ^ ' « « « + i]'«l 0.0055 1 2.167 xlO38 2.167xl038 J

In ß phase region:

a = — î — l n ( ( ;)"""» + [( -)>"■•«» + 1 ] " 2 1 0.0166 1 2.441 x l O 8 2.441 x l O 8 J J

References

[I] Jinyou Wang, Zhiming Ge, and Yanbang Zhou, Titanium Alloy Used in Aerospace Industry. (Shanghai:

Shanghai Metallurgie Press, 1985.)

[2] Rare Metal Materials Processing Handbook [M], Beijing: Metallurgical Industry Press, 1984.

[3]H. Fujii, H. Suzuki, Y. Nakamura. "Transformation characteristics of Ti-6A1-4V alloy on continuous

cooling,"./. Iron Steel Inst. Jpn. 72(5) ( 1986), 299.

[4] T. Seshacharyulu, S.C. Medeiros, W.G. Frazier, et al. "Hot working of commercial Ti-6A1-4V with an

equiaxed a-ß microstructure: materials modeling considerations," Mater. Sei. Eng. A, 263(1999), 272-280.

[5] SX. Semiatin, TR. Bieler. "The effect of alpha platelet thickness on plastic flow during hot working of

Ti-6AI-4V with a transformed microstructure," Ada. Mater. , 49 (2001), 3565-3573.

[6]P. Wanjara, M.Jahazi, H.Monajati, et al., Mater. Sci.Eng. A396(2005)50-60.

[7] CM. Sellars, W.J.McG. Tegart. "On the mechanism of hot deformation," Ada Metallurgica,

14(9)(1966), 1136-1138.

[8] Poirier J P. Delin Guan. High temperature plastic deformation of the crystal ( Dalian: Dalian University

of Technology Press, 1989), 24.

[9] Yongqing Zhao. "Research of the deformation mechanism and the flame retardant property of Ti40

flame retardant alloy"(Ph.D. thesis, Northeastern University, 1998).

[10] Liangyin Xiong, Hongbin Wu, Wei Wang, et al. "Hot deformation of Annealing state TC4 alloy,"

Non-ferrous Metal., 57 (1) (2005), 8-11.

[II] Dyment F. "Self and solute diffusion in titanium and titanium alloys," Titanium'80, Science and

Technology: Proceedings of the Fourth International Conference on Titanium, éd. H Kimura and O Izumi,

(Warrendale Metallurgical Society of AIME, 1980), 510-528.

[12] Aiming Xiong, Shenghui Chen, Weichao Huang, et al. "High temperature deformation of TC6

Titanium alloy and its microstructure development," Rare Metal Materials and Engineering, 32 (6) (2003),

447-450.

[13] Philippart I, Pack H J. "High temperature dynamic yielding in metstable Ti-6.8Mo-4.5F-l.5Al,"

Materials Science and Engineering, A243 (1998), 196-200.

[14] Zener C, Hollomon J H. JApplPhys., 15(1) (1944), 22.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals <£ Materials Society), 2012

EVOLUTION OF MICROSTRUCTURES AND PROPERTIES OF Ti-44Al-6V-3Nb-0.3Y ALLOY AFTER FORGING AND ROLLING

Yuyong Chen*, Hongzhi Niu, Shulong Xiao*, Ping Sun, Changjiang Zhang

National Key Laboratory for Precision Hot Processing of Metals, Harbin Institute of Technology, Harbin 150001, PR China

Abstract

Generally, beta y-TiAl alloys possess excellent hot deformability and good mechanical properties. In this paper, the evolution of microstructures and tensile properties of Ti-44A1-6V- 3Nb-0.3Y (at. %) alloy after hot forging and rolling were investigated systematically. SEM results indicated that the cast microstructure of current TiAl alloy was completely broken down and refined by forging and rolling. By SEM and TEM measurements, the detailed microstructures and phase composition in the as-forged and as-rolled conditions were characterized and compared. Fully lamellar microstructures of the deformed alloy were obtained by heat treatment of 1330°C71h/FC + 900°C710h; Additionally, to assess the effect of microstructure variation on mechanical properties, tensile properties were measured at room temperature and 700 °C. It is an extremely effective mean to optimize the microstructures and enhance the mechanical properties of TiAl alloy by hot forging and rolling. Keywords: Beta y-TiAl alloy, Forging, Rolling, Microstructures, Tensile properties;

Introduction

TiAl based alloys are highly promising for high-temperature structural applications in aerospace and automotive industries, due to their attractive properties, such as low density, good elevated temperature strength, high resistance to oxidation, excellent creep properties [1-2]. However, the low room temperature ductility and poor deformability limit their extensive applications [2-4]. To overcome these problems, tremendous efforts have been devoted to preparing TiAl alloys with refined and homogeneous microstructures, such as alloying [5], heat treatment [6], thermo-mechanical treatment (TMT) [7-10], including isothermal forging, pack forging, hot extrusion and rolling etc. Gamma TiAl sheets have extremely promising aerospace applications, including nozzle tiles for gas turbine engines, nozzles for helicopters, back structures for scramjets, and thermal protection systems for reusable launch vehicles (RLV) etc. [11]. y-TiAl sheets are generally fabricated by ingot-metallurgical (IM) and powder-metallurgical (PM) processing routes [11, 12], Up to now, only the advanced sheet rolling process (ASRP) developed by Plansee AG (Austria), reaches production scale [13]. As to IM process, pack rolling usually proceeds after canned forging steps to ensure the refined microstructure and excellent hot deformability. Understanding the correlations of processing conditions with the resulting microstructures and corresponding

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mechanical properties is of great importance to obtain TiAl alloys with expected service performance. By now, the comparative study of microstructures and properties of TiAl alloys after forging and rolling have been rarely reported comprehensively systematically, especially nearly no relevant reports on beta gamma TiAl alloy sheets have been made. In this study, one beta gamma TiAl alloy sheet fabricated by IM route was introduced. The microstructures and tensile properties of as-cast, as-forged and as-rolled Ti-44Al-6V-3Nb-0.3Y (at. %) alloy were investigated and compared detailedly and systematically.

Experimental procedure

Ti-44Al-6V-3Nb-0.3Y (at. %) sheet was prepared by hot pack rolling after ingot casting and subsequent canned forging process. The TiAl ingot was prepared by induction skull melting (ISM), then was hot isostatic pressed (HIPed) at 1250 °C for 4h under a pressure of 175 MPa, prior to forging. Two-step forging was performed on cylindrical billet (09Ommx 100mm) canned by stainless steel pack at 1200°Cand 0.01s"1, using a total engineering strain of 80%. An 8 mm thick round plate was cut from the mid-depth of the gained TiAl pancake (Fig. 1(a), and then the forged plate was canned by steel jacket and sealed. Crack-free TiAl sheet with dimensions of 360* 170x2.7mm3 was produced through ten-pass hot rolling at 1200°Q with a nominal thickness reduction per pass of approximately 10% and rolling speed approximately 0.5m/s (see Fig. 1(b)). The TiAl samples were held at 1200°C for 1-1.5h before each forging step and the initial rolling pass, as to each subsequent rolling pass, the jacketed TiAl plate was kept at 1200°Cfor 15-30 min. Following the final deformation step or pass, the can-forged pancake or hot-rolled sheet were furnace cooled. In addition, the orientation of the TiAl sheet was invariable during the ten-pass rolling.

Fig. 1. (a) the can-forged Ti-44Al-6V-3Nb-0.3Y alloy pancake, (b) the pack-rolled TiAl sheet fabricated by ingot metallurgy (IM) route (360mmxl70mmx2.7mm).

The microstructures were characterized by OM, SEM and TEM techniques. SEM samples were cut from the center of the pancake or the sheet, and then were prepared by standard procedure. TEM specimens were cut close to the core of pancake (or sheet) with measured surfaces perpendicular to forging direction (or sheet normal), then were prepared by standard mechanical polishing and twin-jet electropolishing. To assess and compare the tensile properties of the present alloy with various microstructures,

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samples cut from the TiAl pancake and sheet were heat-treated in single a phase field at 1320°C for lh to obtain fully lamellar microstructures and aged at 900°C for 10h to stabilize the microstructures. The alpha transus of this alloy was estimated to be 1290°C via heat treatment and metallographic observation. Tensile specimens with gauge section of 20mmx6mmx2mm were cut from the central areas of the cast ingot, forged pancake and the rolled sheet, as well as their heat-treated samples. All the tensile tests were conducted on Instron universal test machine at room temperature with a strain rate of 10"4 s"1.

Results and Discussion

Microstructures before and after hot rolling

Fig. 2 shows Microstructures of as-cast Ti-44Al-6V-3Nb-0.3Y alloy. As shown in Fig. 2(a), current alloy consists of fine lamellar colonies and mixtures of B2 and y phases surrounding the lamellar grains. Straight uVy Lamellae and strip B2 phase in y grain can be clearly seen in Fig.2(b) and (c) respectively.

Fig. 2 Microstructures of as-cast Ti-44Al-6V-3Nb-0.3Y alloy; (a) SEM image in BSE mode, (b) straight oç/y Lamellae and (c) strip B2 phase in y grain existing around lamellar colonies.

Fig.3 shows the deformed microstructures after canned forging and pack rolling. The as-forged alloy displays fully dynamic recrystallization (DRX) microstructure, consisting of y grains (black) and B2 grains (bright) as well as the refined lamellar colonies, Fig. 3(a). The refined lemallar colonies can be clearly observed in the magnified image, Fig. 3(b). Nearly no coarse residual lamellar colonies are observed. As to pack rolling of TiAl sheets via IM route, refined and homogenous DRX microstructure after forging is always popular. Fig. 3(c) shows the typical microstructure of as-rolled beta y-TiAl alloy. The present TiAl sheet reveals fine and uniform microstructure consisting of dominant y (black) and B2 phases (bright). Nearly no lamellar colonies can be observed or detected by SEM observation.

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Fig. 3. Microstructures of Ti-44Al-6V-3Nb-0.3 Y alloy after hot forging and rolling, (a) as-forged, (b) high magnified image of as-forged TiAl alloy; (c) as-rolled.

Fig. 4 shows the TEM images of the deformation microstructures in the forged TiAl pancake and rolled sheet. Fig. 4(a) and (b) represent microstructure of the forged pancake, while Fig. 4(c) and (d) describe that of the rolled sheet. Fragmented lamellar colonies with high density of dislocations can be clearly observed in Fig. 4(a), fine DRX grains in the TiAl pancake are described in Fig. 4(b). Fig. 4(c) represents the decomposition of thin 02 laths and the simultaneous coalescence and effectively coarsening of two neighboring y lamellae, which can not be detected by SEM, as shown in BSE image Fig. 3(c). It's clear that the lamellar spacing is much coarser compared with that in as-forged condition. Noteworthy that the coarsening of y lamellae is induced directly by a lath decomposition (as arrowed). The orientation relationship of y and interrupted 012 lamellae has been confirmed by TEM observation to be the same as the typical crystallographic habit relationship at y/012 planar interface, as illustrated in Fig. 4(d).

Fig. 4. Bright-field TEM images showing deformation microstructures of the current TiAl alloy.

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(a) Fragmented lamellar colonies with high density of dislocations, (b) DRX grains in the as-forged pancake; (c) coarsening of y lamellae by a lath dissolution, (d) corresponding SAED of y/d2 lamellae in current TiAl sheet.

The microstructures of TiAl alloys are generally determined by the alloy nature, thermal processing history and heat treatment. According to the investigations on phase diagram and transformation of high ß phase-contained Ti-Al-V alloys [14] and TiAl-X (X=V, Cr, Mo, W, Ta, etc.) alloys [15], the stable phase composition should be a + y + ß near 1200°C and y + ß (B2) under eutectoid temperature (Te„) for current high ß-stabilizers (i.e. V and Nb) containing TiAl alloy. Consequently, when temperature decreases below Tcu during rolling process, a phase (lamellae) will decompose into y + B2 phase, while y —* a transformation take place when temperature is above Teu (i.e. the heating and holding processes as well as rolling process at temperature above Teu). This is a repeated transformation process. However, in view of transformation thermodynamics and kinetics of current TiAl alloy, only transformation of a lamellae decomposition in prior metastable lamellar colonies caused by forging can be reserved after rolling. The a dissolution was accompanied by growing and coarsening of y lamellae. Finally, the lamellar colonies in as-rolled TiAl sheet are characterized by coarse y lamellae and interrupted 012 laths. By contrast, traditional low ß-stabilizers containing gamma TiAl sheets generally consist of refined gamma grains and a few of lamellar colonies. Such as well-known Gamma Met sheet (Ti-46.5Al-4(Cr, Nb, Ta, B)) and GM PX sheets (TM5Al-5Nb-B-C) reported in Ref. [9], and Ti-45.5Al-2Cr-2Nb sheet reported in Ref. [16], nearly no B2 phase can be found in the as-rolled condition. It is no doubt that the present ß/B2 phase modified TiAl alloy sheet would display excellent hot forming properties.

Fig. 5. Fullly lamellar microstructures of current deformed TiAl alloy after being heat-treated in single a phase field, (a) hot forging + heat treatment of 1330oülh+900°ül0h/FC and (b) hot rolling + heat treatment of 1330oaih+900°<710h/FC.

Generally, via thermal-mechanical treatment, i.e. hot forging, extrusion and pack rolling etc., the beta y-TiAl alloys can be effectively strengthened and their microstructures can be further optimized [5, 7 and 17]. However, fully lamellar TiAl alloys with refined colony size exhibit better comprehensive mechanical properties, especially high temperature creep resistance, than beta phase containing TiAl alloys [18-20], Thanks to heat treatment in single a phase field, the unstable ß phase can be reduced or removed effectively. As shown in Fig. 5(a) and (b), after being heat treated in single a phase field at 1320°C for lh, both current as-forged and as-rolled TiAl specimens display fully lamellar microstructures, with mean grain size of about 80um for

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the former and approximately 150um for the latter. It is clear that the deformed microstructures of the present TiAl alloy grew up dramatically during heat treatment.

Comnarison of tensile properties

The tensile properties of Ti-44Al-6V-3Nb-0.3Y alloy at different processing conditions are listed in Table. 1. We can easily find that after forging or rolling process, the tensile strength and ductility are enhanced significantly both at room temperature and high temperatures ( i.e. 700 °C and 800 *C). After forging, the yield strength (YS) and ultimate strength (UTS) increase from 400MPa and 555 MPa to 620 MPa and 680 MPa respectively, with ductility S enhanced from 0.4 to 0.95% at room temperature (RT); At 700'C, the YS increases from 562MPa to 708 MPa and UTS increases from 650 MPa to 890 MPa after forging, with S increasing to 36%. Current beta y TiAl alloy exhibits obvious "yield stress anomaly" (YSA), i.e. the tensile yield stress increases with the temperature increasing from RT to 700°C By now, the generally accepted mechanism for "YSA" in intermetallic compounds is the thermally activated cross-slip of dislocations and the resulting dislocation locking (Kear-Wilsdorf locks) [21, 22]. While at 800°Q current as-cast alloy processes comparable tensile strength to that after forging, with tensile ductility much lower compared to 95% of the as-forged alloy.

It is obvious that after pack rolling, the tensile strength reduced significantly compared to that of the as-forged condition. The YS decreases from 620MPa of the forged alloy to 520 MPa, and the UTS reduces from 680 MPa to 598 MPa at room temperature, with the RT elongation ( S ) increasing from 0.95% to 1.2%. Meanwhile, at 700 °Q the YS and UTS of the as-forged alloy decreases sharply from 708 MPa and 890 MPa to 570 MPa and 700 MPa respectively after rolling. As well known, mechanical properties of TiAl alloys are highly dependent on their microstructures. It has been universally confirmed that the tensile strength of TiAl alloys can be markedly improved by refining the grain size, increasing the content of lamellar colonies, especially lamellar colonies with refined lamellar spacing [18, 19 and 23]. Meanwhile, lamellar interfaces and a2 laths in the lamellar colonies can be effective barriers to dislocation mobility [1, 24]. In this study, the as-rolled TiAl alloy exhibits larger grain size and coarser lamellar spacing compared with as-forged condition, which is the primary reason for the tensile-strength degradation after rolling. The a decomposition and the decreasing of interfaces can significantly reduce the strength of TiAl sheet. Overall, the tensile-strength degradation caused by hot rolling are caused by coarsening of DRX grains and lamellar spacing, and decrease of a laths.

Table. 1 Tensile properties of Ti-44Al-6V-3Nb-0.3Y alloy at different processing conditions.

As-east

As-forged

As-rolled

Temperature

RT

700°C 800°C RT

700°C 800°C

RT

YS, MPa

400 562 420 620 708 420 520

UTS, MPa

555 650 510 680 890 490 598

Elongation S,%

0.4 7.6 18.5 0.95 36 95 1.2

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Hot forging + 1330°mh+900°CyiOh

Hot rolling + 1330°Clh+900°(710h

700°C RT

700°C RT

700°C

570 630 716 610 755

700 724 870 705 920

31 0.75

15 0.68

12

After heat treatment of 1330°ülh+900o(yi0h, we find that both the tensile strength of the as-forged and as-rolled specimens, with full-lamellar microstructures, are significantly improved, with ductility ( S ) decreasing slightly. The present full-lamellar TiAl alloy features better tensile properties than that of the as-cast alloy. Compared to the YS of 630 MPa and UTS of 724 MPa of the heat-treated forged alloy, the as-rolled TiAl alloy with full-lamellar microstructure exhibits slightly lower tensile strength, with YS decreasing to 610 MPa and UTS reducing to 705 MPa. The removing of the ß phase in the deformed TiAl alloy and the resulting fully lamellar TiAl microstructures significantly contribute to the tensile strength, while the decreasing of the tensile ductility is mainly caused by the coarsening of the grain size during high temperature treatment.

Conclusions

In this study, the microstructure evolution and tensile property variation of Ti-44A1-6V-3Nb-0.3Y alloy were investigated systematically. The main conclusions drawn from this study are as follows. (1) After forging or rolling, the cast microstructure is broken down and refined effectively. The

present deformed TiAl alloy is characterized by completely DRX microstructures, consisting of equiaxed y and B2 grains, as well as some lamellar colonies. The as-forged alloy has more lamellar colonies and fewer B2 grains than the as-rolled alloy. During hot-pack rolling, the DRX grains grow up and the coarsening of y lamellar spacing by a lath decomposition takes place.

(2) Hot forging and rolling processes are extremely effective means to optimize the cast microstructure and enhance the mechanical properties of TiAl alloy. After forging or rolling, the tensile properties of current beta y TiAl alloy are improved significantly. However, the tensile properties of as-forged alloy reduce obviously after hot rolling, with yield strength decreasing from 620MPa to 520 Mpa at room temperature, and from 708 MPa to 570 MPa at 700 °C respectively. The degradation of the tensile strength after hot rolling is primarily ascribed to coarsening of both the DRX grains and lamellar spacing, as well as the decrease of a laths.

(3) After heat treatment in single a phase field, both the forged and rolled TiAl alloys display fully lamellar microstructures, and the corresponding tensile properties are improved further.

References

[1] Y.W. Kim, "Effects of microstructure on the deformation and fracture of y-TiAl alloys," Mater Sei Eng A, 192/193(1995), 519-533.

[2] X.H. Wu, "Review of alloy and process development of TiAl alloys," Intermetallics,

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14(2006), 1114-1122. [3] Y.W. Kim, "Gamma titanium aluminides: their status and future," JOM, 47 (1995), 39-41. 4] Y.W. Kim, "Wrought TiAl alloy design," Trans Nonferrous Met Soc China, 9 (suppl.l)

(1999), 298-308. 5] R.W. Imayev, V.W. Imayev, M. Oehring, F. Appel, "Alloy design concepts for refined

gamma titanium aluminide based alloys," Intermetallics, 15(2007), 451-460. 6] H. Clemens, A. Bartels, S. Bystrzanowski, H. Chladil, H. Leitner. "Grain refinement in

y-TiAl-based alloys by solid state phase transformations," Intermetallics, 14(2006), 1380 -1385.

7] T. Tetsui, K. Shindo, S. Kobayashi, M. Takeyama, "Strengthening a high-strength TiAl alloy by hot-forging," Intermetallics, 11 (2003), 299-306.

[8] T. Tetsui, K. Shindo, S. Kaji, S. Kobayashi and M. Takeyama, "Fabrication of TiAl components by means of hot forging and machining," Intermetallics, 13(2005), 971-978.

[9] S.L. Draper, D. Krause, B. Lerch, I.E. Locci, B. Doehnert, "Development and evaluation of TiAl sheet structures for hypersonic applications," Mater Sei Eng A, 464(2007), 330-342.

[10] T. Carneiro, Y.W. Kim, "Evaluation of ingots and alpha-extrusions of gamma alloys based on Ti-45Al-6Nb," Intermetallics, 13(2005), 1000-1007.

[11] G. Das, H. Kestler, H. Clemens, P.A. Bartolotta, "Sheet gamma TiAl: Status and opportunities," JOM, 56(2004), 42-45.

12] H. lemens, H. Kestler, "Processing and applications of intermetallic y-TiAl-based alloys," Adv Eng Mater, 2(2000), 551-570.

;i3] H. Clemens, H. Kestler, N. Eberhardt and W. Knabl. In: Kim YW, Dimiduk D, Loretto M, editors. Gamma titanium aluminides. Warrendale, PA, USA: TMS; 1999. p. 209-223.

14] M. Takeyama, S. Kobayashi, "Physical metallurgy for wrought gamma titanium aluminides: Microstructure control through phase transformations," Intermetallics, 13(2005), 993-999.

[15] R. Kainuma, Y. Fujita, H. Mitsui, I. Ohnuma, K. Ishida, "Phase equilibria among a (hep), ß (bec) and y (Llo) phases in Ti-Al base ternary alloys," Intermetallics, 8(2000), 855-867.

[16] S.L. Semiatin, B.W. Shanahan, F. Meisenkothen, "Hot rolling of gamma titanium aluminide foil," Acta Mater, 58(2010) 4446-4457.

[17] X.J. Xu, J.P. Lin, Y.L. Wang, J.F. Gao, L. Lin and G.L. Chen, "Effect of forging on microstructure and tensile properties of Ti-45Al-(8-9)Nb-(W, B, Y) alloy," J Alloys Compd, 414(2006), 175-180.

[18] C.T. Liu, J.H. Schneibel, P.J. Maziasz, J.L. Wright, D.S. Easton, "Tensile properties and fracture toughness of TiAl alloys with controlled microstructures," Intermetallics, 4(1996), 429-440.

[19] G.X. Cao, L.F. Fu, J.G. Lin, "Relationships of microstructure and properties of a fully lamellar TiAl alloy," Intermetallics, 8(2000), 647-653.

20] J.G. Wang, T.G. Nieh, "Creep of a beta phase-containing TiAl alloy," Intermetallics, 8(2000), 737-748.

[21] F.R.N. Nabarro, Dislocations in Solids (Elsevier, Amsterdam, 1996), vol. 10. [22] B.A. Grinberg, M.A. Ivanov, "Intermetallic compounds NJ3A1 and TiAl: Microstructure and

Deformation Behavior," Ural Division of the Russian Academy of Sciences, Yekaterinburg, 2002 (in Russian).

[23] R.T. Zheng, Y.G. Zhang, C.Q. Chen, G.A. Cheng, "The ambient temperature tensile behavior of duplex y-TiAl-based alloys," Mater Sei Eng A, 362(2003), 192-199.

[24] F. Appel, R. Wagner, "Microstructure and deformation of two-phase y-titanium aluminides," Mater Sei Eng R, 22(1998), 187-268.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

EFFECT OF FORGING ON MICROSTRUCTURAL CHARACTERISTIC AND TENSILE PROPERTIES OF IN SITU (TiB+TiC)/Ti COMPOSITE

Yuyong Chen*, Changjiang Zhang, Shulong Xiao*, Dezong Wu, Hongzhi Niu

National Key Laboratory of Science and Technology on Precision Heat Processing of Metals, Harbin Institute of Technology Harbin 150001, P. R. China

Keywords: Titanium matrix composites, Microstructure, Mechanical properties, Forging.

Abstract In this work, 2.5vol. % (TiB+TiC)/Ti composite was prepared by in situ casting route then

1-D forging. The microstructure and tensile properties were presented and discussed. The results indicate that the as cast microstructure can be significantly modified by 1-D forging. After forging, TiB and TiC segregated at the prior ß grain boundaries within the as-cast composite tend to fracture and align perpendicular to forging direction. Reduction in aspect ratio of reinforcements and a lath is also observed. 1-D forging can enhance the strength and elongation of as cast composite significantly. However, the increment in strength is quite limited as strain temperature increases to 700 °C. Additionally, Room temperature and high temperature fracture mechanisms are also discussed as well.

1 Introduction Discontinuously reinforced titanium matrix composites (DR-TMCs) offer considerable

potential for improvement in properties and service temperature due to their high strength and elevated temperature resistance and stiffness, which extends the application field such as automotive, aerospace and advanced weapon system [1-3]. Among the reinforcements frequently-used in the titanium alloys such as SiC, TiC, AI2O3, TiB and T1B2, TiB and TiC are considered as the most effective ceramic reinforcements due to the favorable modulus, stability and similarity of heat expansion coefficients [3-5], Additionally, TiB and TiC reinforced DR-TMCs have been extensively demonstrated to possess superior mechanical properties over the single TiB or TiC reinforced DRTMCs [6-7].Nowadays, in situ casting route has been utilized to prepare DR-TMCs due to its obvious advantages such as direct casting forming, simplicity, net-shape, low cost and flexibility[3, 6, 8].

Currently, one drawback associated with as-cast (TiB+TiC)/Ti composites is their low ductility at room temperature (especially at high volume fraction of reinforcement), which limits the broader applications. So, it is necessary to enhance the mechanical properties of as-cast (TiB+TiC)/Ti composites by second processing [9-11], Wang et al [10] obtained relatively high elongation to failure of 8.9% by using hot forging on as-cast 2vol. %(TiB+TiC)/Ti-1100 composites. Ivasishin et al [11] found that forging and rolling in the ß phase field can improve reinforcement distribution and refine matrix microstructure in Ti-6A1-4V-1.55B eutectic alloy to enhance the mechanical properties. However, the second processing in previous studies are complex and high energy-consuming. Therefore, it is quite necessary to employ the one-dimensional forging (1-D forging) technique possessing simple and energy-saving advantage to enhance the mechanical properties of (TiB+TiC)/Ti composites.

In the present study, 2.5vol. % (TiB+TiC)/Ti composite was prepared by in situ casting route then 1-D forging. The main objective is to investigate effect of forging on microstructure

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and tensile properties of (TiB+TiC) /Ti composite. Room temperature and high temperature fracture mechanisms are discussed as well.

2 Experimental The material used in the experiment was a Ti-6Al-2.5Sn-4Zr-0.7Mo-0.3Si-0.2Y matrix alloy

reinforced with 2.5 vol. % (TiB+TiC) (TiB: TiC=l:l). In synthesizing composite, the raw materials were sponge titanium (99.9%), B4C powder (98%), graphite powder (99%) and other alloy elements. The raw materials were blended and then melted in an induction skull melting furnace (ISM). TiB was in situ synthesized by the following chemical reactions:

Ti + C = TiC (1) B4C + 5Ti = 4TiB + TiC (2)

The composite melt was poured into a metal mould to form cylindrical ingot with a length of 120mm and a diameter of 140mm.A cylindrical ingot with a diameter of 80mm and a length of 120mm was used for one-dimensional forging (1-D forging) carried out at 1050 °C and at a strain rate of 0.013s'1 followed by air cooling. The sample was preheated at 1100 °C for 50 min and then 1-D forged with a total deformation of 70%.

Samples for microstructural observation and tensile test were cut by electric discharging from the as cast ingot and 1-D forged pancake. The phase constituents and microstructure of the produced samples were examined by X-ray diffraction (Riguku D/MAX RB, Cu target, 50kv/50mA) and SEM (FEI Quanta 200F). Image analysis software was used to characterize the microstructural parameters and reinforcements. Flat tensile specimens were performed on Instron 5500R testing machine at room temperature, 600 °C and 700 °C, driven at crosshead speed of 0.5 mm«min-l. At least three tests were performed on each condition and average values were reported. Fracture surfaces were observed using Hitachi S-570 SEM.

3 Results and Discussion 3.1 Microstructural characterization

Fig. 1 shows the X-ray diffraction pattern of the as-cast composite. It is shown that Ti, TiB and TiC phases are present and that no other metastable borides and carbides are detected, indicating that (TiB+TiC)/Ti composite can be synthesized through the in situ reaction during induction melting process.

Fig.l X-ray diffraction pattern of as-cast composite The microstructures of as cast (TiB+TiC)/Ti composite are shown in Figure 2. As shown in

Fig.2a, the composite consists of a-ß two-phase matrix and reinforcements decorating prior ß grain boundaries. Two main morphologies of reinforcements, i.e. short-fiber shape and equiaxed or near-equiaxed shape are observed in Fig.2b. Spectrum analysis techniques are used to

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distinguish between the two reinforcements, revealing that the reinforcements with whisker shape are TiB whereas those with equiaxed or near equiaxed shape are TiC. The difference in morphologies between TiB and TiC reinforcements can be mainly related to their difference in crystal structures [12].

Fig.2. Microstructure of as cast (TiB+TiC)/Ti composites (a) low magnification and (b) high magnification.

Fig.2 also shows the significant features include a prior ß grain size of about 138um and an a lath width of about 6.32um. While the prior ß grain size and a lath width of as cast high temperature titanium alloy are approximatively lOOOum and lOum respectively, according to previous studies [13, 14]. Hence, it can be inferred from current results that TiB or TiC reinforcements have a significant refinement in the matrix microstructure within the composite. The most widely accepted mechanism of refinement (especially prior ß grain size) was constitutional supercooling due to the addition of B proposed by Tamirisakandala [13]. fn the current study, the B and C sources of the TiB and TiC reinforcement are B4C and C powder, the in situ reaction (1) and (2) can occur far below the melting point. When the melting temperature exceeds the liquidus temperature, TiB and TiC were completely melted in to the melt. According to the Ti-B and Ti-C binary phase diagram, the compositions of the current materials are hypoeutectic, which means prior ß phase firstly generated followed by the precipitation of TiB and TiC. As the ß phase nucleates, the boron and carbon are rejected into the liquid leading to solute enrichment in the liquid ahead of the solid-liquid interface resulting in constitutional supercooling. Undoubtedly, this phenomenon led to instability in solid-liquid interface and formation of more nuclei. Once nuclei are formed, the presence of boron and carbon rich layer retards the grain growth allowing more grains to form in the surrounding area. Afterwards, the TiB and TiC precipitate at the prior ß grain boundaries, as shown in Fig.2.

Fig.3 shows the microstructures of 1-D forged (TiB+TiC)/Ti composite. It can be seen that reinforcements are randomly distributed in the cross-section (Fig.3a) while they are inclined to align perpendicular to forging direction (Fig.3b). It can be observed that many of the high-aspect ratio TiB whiskers are broken into short segments leading to reduction in aspect ratio and the matrix flow fills the gaps resulting in no noticeable voids. During deformation in the two o + ß phase field, the evolution of prior lamellar a is depended on the orientation relative to the compression axis of the specimen. The prior lamellar a which are perpendicular to the compression axis are elongated and become thinner. Additionally, it is obvious to see from Fig.3b that the secondary a lath is significantly refined after forging. The refinement in the secondary a lath can be attributed to the presence of TiB and relatively high cooling rate (air cooling) after forging. The presence of TiB can enhance the kinetics of phase transformation (ß —» ß + a) by providing additional nucleation sites. Indeed, Bermingham et al [15] have further

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confirmed that TiB assists the nucleation of a phase. At the same time, the relatively high cooling rate can enhance the kinetics of phase transformation as well.

Fig.3. Microstructure of as forged (TiB+TiC)/Ti composite: (a) cross-sections and (b) longitudinal sections 3.2 Tensile properties

Tensile properties of (TiB+TiC)/Ti composite were investigated after 1-D forging at 1050°C The main purpose of hot forging is ingot breakdown and eliminating shrinkage cavities formed during solidification. The tensile properties of the as forged (TiB+TiC)/Ti varying with temperatures during tensile test were shown in Table. 1 and compared with those of the as-cast composite. It can be seen from Table 1 that the room temperature ultimate tensile strength (UTS), yield strength (YS) and elongation to failure (8) of the as cast composite significantly increase after 1-D forging. Meanwhile, reduction in strength and increase in elongation for the composite after forging are noticeable with the rising of temperature. This tendency is more obvious as the tensile temperature rising to 700 °C. This can be due to softening of matrix within the composite at high temperature. Additionally, after 1-D forging, the increment in strength at high temperature is not notable compared with that tested at room temperature.

Table I. Tensile Properties of As Cast and As Forged (TiB+TiC)/Ti Composites

Materials

As cast 1-D forged

Room temperature YS UTS

(MPa) 1039.8 1192.3

(MPa) 975.5 1103.2

8(%)

1.98 4.95

UTS

(MPa) 646.5 691.2

600 °C YS

(MPa) 543.2 567.4

8 (%)

6.6 12.5

UTS

(MPa) 502.9 520.2

700 °C YS

(MPa) 413.2 445.3

8 (%)

17.4 25.6

According to the classic Hall-Petch equation: <ry= ao+ KydT , where ayii the yield strength, ob and Kyare material constants, and d is the mean grain size. The value of Ky generally depends on the number of slip systems and is higher for HCP metals than that for FCC and BCC metals [16]. Since a-Ti is HCP, the grain size affects the yield strength significantly. Lütjering [17] has pointed out that the most influential microstructural parameter on the mechanical properties of fully lamellar structures is the a colony size as it determines the effective slip length. The a lath width of (TiB+TiC)/Ti composite is significantly refined after 1-D forging, as a result the yield strength of the composite increases. Additionally, the refinement in a lath can lead to improvement in elongation as well because the thinner a lath can effectively reduce stress concentration and inactivate the nucleation of crack during tensile test.

Tensile fracture surfaces of the composites with different processing states at room and high temperatures are shown in Fig.4. A brittle cleavage fracture mode was observed in the as-cast composite. The s cracked TiB and TiC can be found on the fracture surface (Fig. 4a). After

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forging, the fracture of the composite presents some characteristics of ductile fracture due to some dimples observed (Fig.4 d). This phenomenon indicates that the fracture type of as cast composite changes from brittle fracture mode to a combination of brittle and soft mechanisms.

While at high temperature, more dimples and less cracked TiB and TiC are observed in the fracture surfaces of as forged composite (Fig.4e) compared with those of as cast composite (Fig.4b), indicating that the fracture type is soft mechanism. This is consistent with the higher elongation to failure of the as forged composite. Additionally, few reinforcements pulled-out, especially at 700 CC, can be seen in Fig. 4f, showing that a large amount of plastic deformation of the matrix exceeds the reinforcements. The phenomenon indicates the additional load carried by reinforcements is quite limited at 700 °C.

Fig.4 SEM images of fracture surfaces of (TiB+TiC)/Ti composite with different processing states at different temperatures: (a) as cast, room temperature, (b) as cast, 600 °C, (c) as cast, 700 °C, (d) as forged, room temperature, (e) as forged, 600 °C, (f) as forged, 700 °C.

Conclusions

1. As-cast microstructures of (TiB+TiC)/Ti composite are significantly modified by 1-D forged. After forging, TiB and TiC segregated at the prior ß grain boundaries within the as-cast composite tend to fracture and align perpendicular to forging direction. Reduction in aspect ratio of reinforcements and a lath is also observed.

2. 1-D forging can enhance the strength and elongation of as cast composite significantly. However, the increment in strength is quite limited as strain temperature increases to 700 °C.

3. Room temperature fracture mechanism of as-cast composite is a brittle cleavage fracture

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mechanism, while the fracture mechanism of the as-forged composite is a combination of brittle and soft mechanisms, soft fracture mechanism dominates the fracture modes at high temperature for both the as cast and as forged composites

References

1. T. Godfrey, P.S. Goodwin, and CM. Ward-Close, "Titanium Particulate Metal Matrix Composites Reinforcement, Methods, and Mechanical Properties," Adv Eng Mater, 2 (2000), 85-92.

2. C.H. Weber, Z.Z. Du, and F.W. Zoky, "High Temperature Deformation and Fracture of a Fiber Reinforced Titanium Matrix Composite," Acta Metall Mater, 44 (1996), 683-695.

3. S.C. Tjong, Y.W Mai, "Processing-Structure-Property Aspects of Particulate-and Whisker-Reinforced Titanium Matrix Composites," Compos Sei Technol, 68 (2008), 583-560.

4. L.J. Huang, L.Geng, H.X. Peng, "In Situ (TiBw+TiCp)/Ti6A14V Composites With a Network Reinforcement Distribution," Mater Sei Eng A, 527 (2010), 6723-6727.

5. CF. Yolton, "The Pre-alloyed Powder Metallurgy of Titanium with Boron and Carbon Additions," X>M, 56 (5) 2004, 56-59.

6. W.J. Lu, et al., "Microstructure and Tensile Properties of In Situ (TiB+TiC)/Ti6242 (TiB:TiC=l:l) Composites Prepared by Common Casting Technique," Mater Sei Eng A 311 (2001), 142-150.

7. B.V. Radhakrishna, J. Subramanyam, V.V. Bhanu, "Preparation of Ti-TiB-TiC & Ti-TiB Composites by In-Situ Reaction Hot Pressing," Mater Sei Eng A, 325 (2002), 126-130.

8. S.C. Tjong, Z.Y. Ma, "Microstructural and Mechanical Characteristics of In Situ Metal Matrix Composites," Mater Sei Eng R, 29 (2000), 49-113.

9. H.T. Tsang, CG. Chao, C.Y. Ma, "Effects of Volume Fraction of Reinforcement on Tensile and Creep Properties of In-Situ TiB/Ti MMC," Scripta Mater, 37 (1997), 1359-1365.

10. M.M. Wang, et al., "Effect of Volume Fraction of Reinforcement on Room Temperature Tensile Property of In Situ (TiB+TiC)/Ti Matrix Composites," Mater Des, 27 (2006), 494-498.

l l .O.M. Ivasishin, et al., "Processing, Microstructure, Texture, and Tensile Properties of the Ti-6A1-4V-1.55B Eutectic Alloy," Metall Mater Trans A, 39 (2008), 402-416.

12. X.N. Zhang, W.J. Lu, D. Zhang, "In Situ Technique for Synthesizing (TiB+TiC)/Ti Composites," Scripta Mater, 41 (1999), 39-46.

13. S. Tamirisakandala, et al., "Grain Refinement of Cast Titanium Alloys via Trace Boron Addition," Scripta Mater, 53 (2005), 1421-1426.

14. V.K. Chandravanshi, et al., "Effect of Minor Additions of Boron on Microstructure and Mechanical Properties of As-cast Near a Titanium Alloy," Metall Mater Trans A, 41 (2010), 936-946.

15. M.J. Bermingham, et al., "Effects of Boron on Microstructure in Cast Titanium Alloys," Scripta Mater, 59 (2008), 538-541.

16. R. Armstrong, et al., "The Plastic Deformation of Polycrystalline Aggregates," Phil Mag, 1 (1962), 45-58.

17. G. Liitjering, "Influence of Processing on Microstructure and Mechanical Properties of (a+ß) Titanium Alloys," Mater Sei Eng A, 243 (1998), 32-45.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Microstructure and Mechanical Properties of Ti-6A1-4V Fabricated by Selective Laser Melting

Marco Simonelli1, Yau Yau Tse1, Christopher Tuck2

'Department of Materials, Loughborough University 2School of Mechanical and Manufacturing Engineering, Loughborough University

Keywords: Selective Laser Melting; Ti-6A1-4V microstructure; martensitic a' phase;

Abstract

Selective laser melting has been recently proposed as a new powder bed manufacturing technology for Ti alloys. In this research, the 3D microstructure characterization of the Ti-6A1-4V pre-alloyed plasma atomized powder material was carried out. Cubic components were then built from the same powder directly onto the build plate or on top of secondary supporting structures using a modulated pulsed fibre laser. The microstructure and mechanical property of the components were analyzed. Spherical and flat pores were found, with supported components showing a larger average porosity probably due to a slower experienced cooling rate. The microstructure of the parts consists mainly of acicular a' needles with only a small volume fraction of retained ß phase. Transmission electron microscopy (TEM) and electron backscattered diffraction (EBSD) studies confirmed the crystallographic nature and the orientation of the a' needles, respectively. The a' needles are mainly arranged in a colony pattern, but basketweave morphology is visible in the bottom region of the supported components. Columnar grains extending across deposited layers were observed as a result of the several thermal cycles which are typical in selective laser melting.

Introduction

New additive manufacturing technologies of titanium alloys have received increasing attention, particularly selective laser melting (SLM) [1,2]. SLM enables the fabrication of parts of any shape directly from a computer aided design model in a single step. It has been reported that the typical microstructure of the SLM processed components consists of oriented columnar grains, intra-layer porosity, and fine acicular laths [3-6]. By varying the orientation of the components on the building platform, it was demonstrated that it is possible to control the size and growth direction of the defects and the columnar grains of Ti-6Al-7Nb components [7], AlNbTi2 precipitates located inside and at the boundaries of martensite plates were also found [7]. Parallel bands corresponding to the Ti3Al phase which would have precipitated during successive deposition steps were observed on Ti-6A1-4V parts [8]. The volumetric energy density applied to the powder bed plays an important role in controlling the porosity alignment and their average size [8]. SLM Ti-6A1-4V is mainly constituted of a' phase; the acicular needles martensitic phase resulted from high cooling rate which is typical during SLM process [9], The microstructure of SLM Ti-6A1-4V is, however, still a matter of discussion as it presents several features that have not been observed in traditionally manufactured components [9], In fact, a clear understanding of the as-built microstructure could provide a valuable basis to optimize the process parameters and achieve products with high mechanical properties. The aim of this research is to compare the microstructure of SLM Ti-6A1-4V components built directly onto the

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building platform or on top of secondary supporting structures. The microstructure of the starting powder material was also characterized in 3D to study the true geometrical shape of the a laths.

Materials And Methods

The starting material is a pre-alloyed plasma atomized Ti-6A1-4V grade 23 with nominal chemical composition as shown in Table I.

Table I. Chemical composition of Ti-6A1-4V powders

Element wt%

N

0.03 C

0.08

H 0.01

Fe 0.02

O 0.13

Al 6.5

V 4.5

Ti Bal

The powder was supplied by LPW, UK. The powder size and distribution were determined using a Malvern MasterSizer machine. Simple cubic components of volume of 1 cm3 were built using an SLM MTT 250. The machine is equipped with a fibre modulated pulse laser, with a maximum power of 200 W. The wavelength of the laser is 1070 nm. The adopted processing parameters are summarized in Table II.

Laser Power [W] Hatch Spacing [urn]

Layer Thickness [um] Point Distance [urn] Exposure Time [us]

Energy Density [J mm" ]

200 180 50 50 251

11.1 * 10"5

Table II. Process Parameters Figure 1. Top and cross-sectional view of the secondary supporting structures

A multi-directional meander scan strategy was used, i.e. the laser scan direction was rotated by 67° at each layer. Seven cubes were built directly on top of a Ti- 6A1-4V build plate (p-samples) and onto Ti-6A1-4V needle-shaped supports (s-samples). The arrangement of the secondary structures of the s-samples is shown in Figure 1. Different support mechanisms provided different contact area and hence different thermal histories to the built component. The microstructural evolution of the SLM built components on different heat sinks were studied in three orthogonal planes as shown in Figure 2.

Figure 2. Schematic diagram showing the investigated planes

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All samples surfaces were prepared by standard metallographic preparation procedure and then etched by Kroll's solution for optical/SEM studies or colloidal silica polished for EBSD mapping. Optical microscopy was carried out using a Reichert-Jung MEF3 optical microscope. The microstructural morphology and orientation were determined using SEM and EBSD. In order to gain further understanding of the fine microstructure, thin TEM disks were prepared. Electropolishing was carried out in a temperature range of -20 to -30°C at 20V with the A2 electrolyte supplied by Struers. The TEM analyses were performed on a JEOL 2000FX TEM operated at 200 kV. The 3D microstructure characterization was carried out on a FEI Nova 600 Nanolab Dualbeam FIB/FEG-SEM. The serial sectioning, stage rotation and EBSD acquisition were performed automatically. A total of 200 slices of 200 nm each were milled, analyzed and 3D reconstructed using Avizo® software.

Results And Discussion

The measured particle size distribution was within 15 and 70 urn, but, about 75% of the examined particle size ranged between 25 to 50 urn. Particles appear spherical, smooth with few smaller condensed particles attached to bigger particles; similar results have been reported elsewhere [10]. When the FIB-milled cross section was studied, a fully lamellar microstructure is found, as indicated in Figure 3a. The EBSD result shows the lamellae are mainly a phase and there is only a limited amount of retained ß phase is present (less than 4%). This may be due to the high cooling rate experienced by the powder during plasma atomization (PA). The orientation of the lamellae is shown in Figure 3b. The examined particle exhibits only few orientation components.

Figure 3. a) EBSD image quality map of a randomly selected particle and b) Inverse pole figure orientation map of a particle. The inset shows the inverse pole figure colour scheme.

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Figure 4. 3D-reconstruction of few a laths

The focused ion beam sectioning series show the PA powder is fully dense [10]. The 3D representation of the a laths in the particles is shown in Figure 4. Each colour represents a particular crystal orientation component, a plates grow throughout the particle and no obvious basket-weave structures are observed in the particle. A few smaller lamellae grow almost perpendicular to larger lamellae. All the built cubes (p- and s-samples) show good geometrical accuracy. When investigating the frontal and lateral surface of the samples, two types of pores are found: pores are flat and oriented in the scanning plane or spherical (pin-hole porosity). S-samples show however a larger average porosity. Although the built cubes show good geometrical accuracy, balling of the melt pool could have arisen because of either too low or too high volumetric energy density during the process [11]. This could have led to the formation of the pores. Porosity might also origin from the presence of hydrogen in the powder. At high temperatures, above the melting point, H would have a natural tendency to bubble out from the surface of the melt pool leaving behind spherical pores [12]. Although p- and s-samples show similar shape of the pores, s-samples show a larger average porosity than p-samples as shown in Figure 5a and b. Even if all the samples were processed using identical parameters (i.e. volumetric energy density) it must be noted that heat conduction of the supported samples is worse than those built directly onto the plate. In fact, in s-samples heat is preferentially conducted through the supports. As the plate is located below the supports it probably did not play a crucial role in the cooling down of the parts. On the other hand, p-samples used the base plate as heat sink which can quickly draw away the heat generated by the scanning laser. As a result of slower heat conduction, the melt pool lifetime in s-samples was extended favoring Marangoni effect to take place and therefore larger pores resulted. The long lifetime of the melt pool would also allow more time for H to coalescence and grow in larger bubbles leaving behind larger spherical pores.

Columnar grains are observed in both p- and s- samples and they grow across several layers. During the laser scan, the tops of the grains in the previous layer can be partially re-melted and they can undergo epitaxial growth in the next layer generating the columnar grains [13]. The columnar grains do not possess a preferential tilt with respect to the building direction: this is probably due to the adopted scanning strategy (the scan direction is rotated by 67° at each layer) [14]. In this current study, the average width of the grain is 210 |am (SD 50 um), which is similar to the scan track width and this may related to the adopted scan strategy [14].

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Figure 5. Porosity on a lateral cross-section of a) p-sample, b) s-sample

Figures 6a and b show a comparison between the bottom region of the supported and plate components. EBSD results show that the acicular needles have an HCP crystal structure. Both the diagrams show that parallel needles have identical orientation. The s-sample displays a more complex entanglement of the acicular needles and a worse image quality map, which implies the presence of higher dislocation density and higher residual stresses (Figure 6a).

Figure 6. [100]/RD Inverse pole figure map of the bottom region on the lateral surface of a) s-sample, b) p-sample; the inset shows the inverse pole figure colour scheme.

It is clearly shown in Figure 6b that p-samples have better image quality. Although p-samples are cooled down more rapidly, heat flow might be more uniform due to the direct contact of the samples to the building plate, leading to a microstructure with less/uniform residual stress. However, further analysis, such as EBSD scans near the top regions and TEM dislocation density analysis, are required to understand this difference. As confirmed by optical analysis, in the s-samples the arrangement of these acicular needles varies according to their location; near the external surfaces (wall surfaces) and from the middle to the top region, the morphology exhibits a colony pattern with parallel acicular needles (Figure 7a). Most of the acicular needles lie at ±

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45° to the building direction. In the bottom region and especially in the component area between the secondary supporting structures, the morphology of the needles shows a basket-weave pattern (Figure 7b). Some of the acicular needles appear entangled as a result of competitive growth. This morphology has been reported in wire deposited Ti-6A1-4V specimens [15]. This basket-weave morphology does not appear on p-samples, where the acicular needles appear mainly parallel one to another. As basket-weave patterns are associated with slower cooling rates it could be supposed that this microstructural morphology forms only in areas in contact with the loose powder which implies a less effective heat sink.

Figure 7. a)Typical wall surface and top region microstructural morphology, b) basketweave morphology, found prevalently in the bottom portion of the specimens

Figure 8. a) Bright field TEM micrograph showing several acicular a' needles, b) indexed SADP with beam direction « [12 13]a-.

The internal morphology of the acicular needles is shown in Figure 8a. Most of the needles appear as long, narrow and parallel laths. This is the morphology attributed to the a' martensitic laths [16]. The corresponding indexed selected area diffraction (SAD) pattern of the marked area in Figure 8a is shown in Figure 8b. Tilting series has confirmed that the martensitic acicular needles possess an hep structure which is very similar to the a equilibrium phase [17]. These needles normally contain high dislocation density, stacking faults and twins, which are the typical substructures of the a' martensitic phase of titanium [16]. The a' martensitic nature of the found acicular needles was further confirmed by quantitative TEM-EDX analysis. The

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needles show a higher content in V, 5.9±0.8 wt%, typical of a'[18]. The tensile properties of the as-built components tested in different directions are summarized in Table III. The as-built components display a certain degree of anisotropy due to the additive layer-wise nature of the process. The presence of pores and fine a' martensitic phase may account for the poor ductility of built parts.

Table III. Results of mechanical tests on the as-built components

Conclusions

In this study, SLM Ti-6A1-4V cubic parts were fabricated from pre-alloyed plasma atomized powders. The components were built directly onto a Ti-6A1-4V build plate or on top of Ti-6A1-4V secondary supporting structures. The plasma atomized powder material exhibits a fully dense microstructure constituted mainly of lamellar a phase. A 3D reconstruction of these a laths has revealed their shape and orientation. Although the aim of this project was to fabricate fully dense components, spherical and flat pores were found in the parts; the components built directly onto the build plate, however, show a smaller average porosity than the supported components. This result is explained considering a probable different heat loss condition experienced by the components. The components present columnar grains that grow across the processed layers along the build direction. The columnar grains are composed of acicular a' martensitic phase. These a' acicular needles appear mainly parallel one to another, but in the bottom region of the supported components a basket-weave pattern can be distinguished. The poor ductility of the built parts is due to the existence of pores and fine a' phase.

Acknowledgments

The authors acknowledge Dr Chang Jing Kong for the tensile testes.

References

1. Ian Gibson, "Additive manufacturing technologies: Rapid prototyping to direct digital Manufacturing", 2009.

2. Yadroitsev I et al., "Strategy of manufacturing components with designed internal structure by selective laser melting of metallic powder". Applied Surface Science 2007;254:980-983.

3. Zhang L et al., "Manufacture by selective laser melting and mechanical behavior of a biomédical Ti-24Zr-4Nb-7.9 Sn alloy". Scr Mater 2011.

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4. Abd Aziz I. "Microstracture and Mechanical Properties of Ti-6A1-4V Produced by Selective Laser Sintering of Pre-alloyed Powders". 2010.

5. Santos E et al., "Microstracture and mechanical properties of pure titanium models fabricated by selective laser melting". Proc Inst Mech Eng Part C 2004;218:711-719.

6. Sercombe T et al., "Heat treatment of Ti-6Al-7Nb components produced by selective laser melting". Rapid Prototyping Journal 2008;14:300-304.

7. Chlebus E, et al., "Microstracture and mechanical behaviour of Ti-6Al-7Nb alloy produced by selective laser melting". Mater Charact 2011.

8. Thijs L, et al., "A study of the microstructural evolution during selective laser melting of Ti-6A1-4V". Acta Materialia 2010;58:3303-3312.

9. Facchini L, et al., "Ductility of a Ti-6 Al-4 V alloy produced by selective laser melting of prealloyed powders". Rapid Prototyping Journal 2010;16:450-459.

10. Ahsan MN et al., "A comparative study of laser direct metal deposition characteristics using gas and plasma-atomized Ti-6Al-4 V powders". Materials Science and Engineering: A 2011.

11. Gu D, et al., "Balling phenomena in direct laser sintering of stainless steel powder: Metallurgical mechanisms and control methods". Materials & Design 2009;30:2903-2910.

12. Mohandas T et al., "Microstracture and mechanical properties of friction welds of an [alpha] [beta] titanium alloy". Materials Science and Engineering A 2000;289:70-82.

13. Wu X, et al., "Microstructures of laser-deposited Ti-6A1-4V". Mater Des 2004;25:137-144.

14. Thijs L, et al., "A study of the microstructural evolution during selective laser melting of Ti-6A1-4V". Acta Materialia 2010;58:3303-3312.

15. Baufeld B, et al., "Wire Based Additive Layer Manufacturing: Comparison of Microstracture and Mechanical Properties of Ti-6Al-4 V Components fabricated by Laser-beam Deposition and Shaped Metal Deposition". J Mater Process Technol 2011.

16. Ahmed T, et al., "Phase transformations during cooling in [alpha] [beta] titanium alloys". Materials Science and Engineering A 1998;243:206-211.

17. Leyens C "Titanium and titanium alloys". : Wiley Online Library, 2003.

18. Qazi J, et al., "Phase transformations in Ti-6Al-4V-x H alloys". Metallurgical and Materials Transactions A 2001;32:2453-2463.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

Computational Modeling of the Dissolution of Alloying Elements

Jun Ou, Aniruddha Chatterjee, Carl Reilly, Daan M. Maijer, Steve L. Cockcroft

The Department of Materials Engineering, The University of British Columbia

Vancouver, BC, V6T 1Z4, Canada

Keywords: Computational Modeling, Dissolution, Ethanol-Water Analogue System, Electron Beam Melting, Ti-6A1-4V

Abstract

The dissolution of alloying elements within liquid titanium alloys is of utmost importance in ingot production to ensure chemical homogeneity and final downstream component performance. A coupled thermal-composition-fluid flow model based on a commercial Computational Fluid Dynamics package ANSYS-CFX has been developed to describe dissolution and homogenization of a solid alloying element introduced into the bulk liquid. The transport of heat, mass and momentum via diffusion, buoyancy forces, and surface tension forces has been incorporated into the computational model together with a Darcy-based dampening term applied to the fluid transport within the liquid/solid two-phase region. Experimental data from an ethanol-water analogue physical model has been used to validate the numerical model, using a combination of fluid flow pattern and interface shape, and temperature data obtained at discrete locations. The comparison of results between the computational model and experiments shows good agreement, providing some confidence that the model can be used to investigate the dissolution and homogenization processes in the Ti-6A1-4V system.

Introduction

Ti-6A1-4V (Ti64) is one of the most important titanium alloys, which is widely used in many aerospace, automotive, marine and medical applications [1, 2]. Electron beam cold hearth remelting (EBCHR) has emerged over the last decade as a critical technology to produce Ti64 ingot [3]. In the process of melting and casting Ti64 ingots, it is critical that the process guarantees that Al-rich regions do not carry over from liquid metal processing to the cast ingot to avoid the so-called Type II hard alpha defects. These form when aluminum rich material enters the melt and is not adequately dissolved and homogenized before solidification [2]. Simply extending the time that the alloy is molten is not an appropriate solution for this problem, as it will cause excessive evaporation of some high-vapor pressure alloying elements within the vacuum environment. Therefore understanding the dissolution and homogenization of alloying elements, such as aluminum, in titanium during the EBCHR process is critical to producing high quality ingots.

The dissolution and homogenization of material in the EBCHR process involves the transport of heat, mass and momentum. A number of experimental and modeling investigations have been undertaken to investigate these transport phenomena within the EBCHR process [4-10], Several of the models have been shown to make realistic predictions[7, 9, 10]. However, because of the significant challenges associated with conducting experiments on liquid titanium within the EBCHR process, which include high temperatures, the high chemical reactivity of liquid titanium and the vacuum environment, there is little published work that is directly relevant.

The aim of this work is to develop a coupled thermal-composition-fluid flow model to simulate the dissolution and homogenization of aluminum, an important alloying element, in bulk liquid

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titanium. This model must incorporate the significant physical phenomena occurring in the EBCHR process so as to make accurate and realistic predictions. At this preliminary stage in the work, a physical model based on an ethanol-water analogue system has been developed. There are several reasons to use the ethanol-water system as an analogue for solid aluminum in liquid titanium. Firstly, as both ethanol and water are transparent and are in the appropriate physical states at reasonable temperatures, it is possible to directly observe and easily measure the relevant phenomena. Secondly, the solute (ethanol) has a lower melting point than the solvent (water) allowing solid ethanol to be introduced into a higher temperature liquid, which is analogous to solid Al added to liquid Ti. Thirdly, ethanol has a lower density and fourthly, water has a higher surface tension, again both analogous to Al in Ti.

The ethanol-water system will be used to investigate various factors impacting on the dissolution kinetics of a lower melting point solid in a liquid of different composition and to generate data suitable for validation of the computational CFD model. This data is crucial to ensure that the model both includes the relevant phenomena and is capable of accurately describing the dissolution and homogenization process. Once validated, this model will be further developed and validated on the Al/Ti64 system.

Experimental Method

Fig. 1 shows a schematic diagram of the experimental apparatus. The apparatus consists of a transparent cylindrical glass cell containing water (the solvent) into which a solid sample (the solute) is placed. T-type thermocouples were used to measure the temperature at two locations: one, within solute at the centerline (TI) and a second, near the wall of the cell, 5mm from the top surface (T2) as shown in Fig. 1. To examine the effect of varying the composition difference between the solute and solvent a variety of compositions were used for the solid "solute" ranging from pure water (no composition difference) to pure ethanol. In each case, the "solute" was produced by solidifying it in a mold placed in liquid nitrogen.

Since, both solvent and solute in the study are transparent, dyes were used to differentiate them. An organic-soluble dye (insoluble in water) was used to distinguish ethanol from water whereas a water-soluble dye (insoluble in organic liquids) was used for solute in cases when water was both the solute and solvent. The use of dyes allowed the solid/liquid interface to be delineated and provides a means of mapping the composition (mass fraction) profile for investigating mass transport phenomena. The later required additional experiments to develop a relationship between composition and video gray scale value. MATLAB was then used to translate the gray scale appearing on the video images recorded from experiments to mass fraction.

The experimental methodology was as follows: the cylindrical cell was filled to a height of 60 mm with distilled water (solvent) at 45-50 °C. The solute sample with an embedded thermocouple and support (made from wood) was then prepared by solidifying in liquid nitrogen. Once solidified, the solute sample was removed and immediately inserted in the solvent. Temperature and video data were then recorded using a data acquisition system. In the case of pure water as the solute, due to expansion of water during solidification, it was necessary to first heat the solid sample within the mold to close to its melting point in order to remove the sample from the mold.

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Fig 1 Schematic diagram of experimental apparatus

Model Description

A computational model has been developed to simulate the dissolution process using the commercial CFD software package ANSYS CFX V12.1. To allow two different components (solute and solvent) to be modeled, a multi-component model has been used. To capture the relevant physical phenomena, the model has been developed to include thermal and mass transport by diffusion and advection, thermal and compositional buoyancy, thermal and compositional Marangoni forces and Darcy damping within the two-phase (solid/liquid) region. The implementation of some of these features required the development of user-written subroutines for use within the code (this is a capability supported by ANSYS-CFX). The dimensions of the model domain and the initial conditions applied correspond with the experiment.

For the sake of brevity, the governing equations which are solved for temperature, pressure, composition and fluid velocity are not shown here. Readers are referred to [11] for details. The key information regarding the model implementation is presented below.

Assumptions - In the model, the following basic assumptions have been made. Firstly, the free surface of the solvent is flat. Secondly it is assumed that the fluid flow is laminar and the fluid is incompressible, and finally, it is assumed that the heat transfer between the fluid and glass (solvent container) as well as the atmosphere is negligible.

Material Properties - For the multi-component model, the fluid is comprised of two components, solute and solvent. The fluid properties are calculated by the following ideal mixture rule:

l = Çsu+ÇsL ( 1 )

P Psu Psv

where P is the relevant fluid/material property, Psu and Psv are the properties of solute and solvent respectively, and, Csu and Csv are the mass fractions of the solute and solvent, respectively, the sum of which is constraint to be 1.

As the experiments involve a change of phase in the solute from solid to liquid, the latent heat of melting needs to be accounted for in the heat balance. For both ice and ethanol a modified specific heat term is used to account for the latent heat. Since the absorption of this heat over a

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narrow temperature range causes numerical instabilities, water was assumed to melt over a 5K temperature range and ethanol over a 10K range. For water, melting is assumed to occur between 268 and 273 K with a latent heat of 334 kJ/kg. For ethanol, melting is assumed to occur between 149 and 159K with a latent heat of 119 kJ/kg.

In addition to latent heat, the transition from solid to liquid also involves a change in viscosity (note: CFX treats the solid and liquid as part of the same analysis domain. Flow in the solid is suppressed by altering both the viscosity and permeability of the solid material). In contrast, to the approach used for latent heat, for the model base case, the viscosity is assumed to change from 10 to 0.0004 Pa-s over a IK temperature range, representing the change from solid to liquid.

Buoyancy - The baseline buoyancy formulation provided in ANSYS-CFX does not allow compositionally dependent buoyancy. Therefore, in the current model both thermal and compositional buoyancy are accounted for by adding a source term to the momentum equation, as shown in equation (3).

Sutfiuoyancy ~ (ß ~ Pref)9 (3)

where SM (kg m"2 s"2) is the momentum source, g (m s"2) is the acceleration due to gravity, pref

(kg m"3) is a reference density, and p is the current local density which varies with both temperature and composition.

Marangoni Force - A force imbalance occurs when there is a gradient in surface tension caused by thermal or compositional variations. The net force, or Marangoni force as it is called, drives liquid from a region of low surface tension toward a region of high surface tension. A surface tension gradient can be caused by a compositional and/or a temperature gradient[12]. In the model, a user subroutine describing the Marangoni force, which is expressed mathematically by equation (4), is used to apply a shear stress at the free surface of the liquid pool.

f fr r\ oadT dadc ...

where a (N m"1) is the surface tension, T (K) is the temperature, C is the compositional fraction, and n (m) is the directional normal parallel to the surface. The first term -2-— refers to the temperature-gradient-based Marangoni force and the second term -^— refers to the compositional-gradient-based Marangoni force. The terms T-(N m"1 K"1) and T- (N m"1 kg"1) are referred to as the surface tension coefficients with respect to temperature and composition, respectively.

Darcy Damping - The dissolution problem necessarily involves a phase transformation and the development of a two-phase region between the solid solute and liquid solvent. Darcy's law has been applied to mathematically describe the flow damping in the two-phase zone and in the solid. In this model, a second momentum source SMDarcy (kg m"2 s"2) is calculated using equation (5) and added to the momentum balance: [14]

$M,Darcy — ~ ~^u (5)

where ß (kg m"1 s"1) is the viscosity, K (m2) is the permeability. In the absence of experimental permeability data, the Kozeny-Carman expression is used to calculate the permeability[14]:

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K = ^ f (6)

where D (m"2) is a coefficient related to the specific solid/liquid interfacial area in the mushy zone, and fs is the fraction of solid. For simplicity, D is assumed to be a constant, in this study the value of 1.67el0 m"2 was used[14]. The temperature range over which the phase change from solid to liquid occurs - i.e.fs changes from 0 to 1 in Eq. 6 - is set to IK in the model base case, consistent with the range used for the change in viscosity.

Results and Discussion

Two experiments were chosen to validate the model: i) solid water (ice) solute in a water solvent; and ii) solid pure ethanol solute in a water solvent. The former allows the effects of thermal buoyancy and thermal Marangoni forces to be evaluated in the absence of compositional effects, while the latter allows both thermal and compositional buoyancy and Marangoni forces to be examined.

Fluid Flow and Interface Comparison - Starting first with the ice-water case, Fig. 2 shows the model results and the experimental results 10 s after the ice is dipped into the water. As can be seen there is relatively good agreement between the two both in terms of the shape of the solid/liquid interface and the direction of the flow pattern. It should be noted that the width of the solid solute displayed in the figure appears -10% wider than the model due to refraction. The main driver for flow is the drop in water temperature adjacent to the ice, which causes both an increase in density, resulting in a downward flow adjacent to the ice, and a reduction in surface tension on the surface adjacent to the interface, resulting in an inward flow at the surface. The latter increases both heat and mass transfer at the top surface resulting in an increased rate of dissolution at the surface as observed in Fig 2(a) and (b). Note also that the profiles of the interface down the vertical edge of the sample and along the bottom observed in the experiment are also predicted by the model - e.g. the enhanced erosion on the vertical face toward the bottom of the sample and the enhanced erosion on the bottom face toward the center of the sample leading to a concave profile.

Fig 2 Comparison of solid/liquid interface and fluid flow pattern of ice/water case at 10 s (a) Model (vector plotting of velocity); (b) Experiment

Fig. 3 shows the flow pattern of the ethanol/water case at 10 s. In this case it can be seen that the compositional buoyancy associated with the increased concentration of ethanol at the interface dominates over the thermal buoyancy driving the fluid flow upwards along the vertical interface

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of the solid. Also in contrast to the previous case, the compositional Marangoni force at the top surface drives the fluid outward, as the surface tension of the fluid is lower at the interface in association with the locally increased ethanol concentration. Thus the two sets of forces, thermal and compositional driven, are opposing one another. It can be seen that, as a result, the erosion at the top surface of the solute appears reduced in comparison to the ice water case. Note: a dotted line has been used to estimate the profile of the solid adjacent to the top surface, as it is difficult to delineate this shape due to the large concentration of die present in the area. Finally, it is also clear from the comparison between Figs. 3 (a) and (b) that the profile of the bottom interface has shifted from being concave to convex.

Fig 3 Comparison of interface profile and fluid flow pattern of ethanol/water case at 10 s (a) model results and (b) Experiment results

Temperature Evolution - The evolution of temperature with time measured by the two thermocouples is compared with the numerical model results in Fig. 4. Fig. 4(a) presents the results for the ice/water case and Fig. 4(b) the results for the ethanol/water case. The temperature range over which the material melts and latent heat is evolved in the numerical analysis appears in the figures as horizontal dashed lines. As shown in Fig. 4(a), the ice/water case, the mold was warmed up in order to remove the ice from the mold (see description of the experimental procedure) and hence TCI started with an initial temperature of 268 K, which was also adopted as the initial condition for the ice in the model. In contrast, for the ethanol/water case, Fig. 4(b) the initial condition applied in the model for the ethanol was approximately 70K, consistent with the thermocouple data TCI.

Focusing primarily on TCI, it can be seen that the temperature evolution predicted by the model generally correlates well with the experimental measurements in both cases with a few notable exceptions. In the ice/water case the largest discrepancy occurs between 0s and roughly 80s in the interval in which the temperature is approximately constant. This period of little temperature change is associated with the absorption of heat in the solute and ends with the melting of the material in proximity to the thermocouple. The deviation is likely associated with the temperature range over which the latent heat is absorbed in the model, which is between 268 and 273K, as opposed to in the actual case where it is absorbed at 273K in the ice-to-water transformation. In contrast to the ice/water experiment, the ethanol water experiment shows a period of gradually increasing temperature between 0 and 25s, at which point the ethanol melts. Also there does not appear to be a substantial plateau associated with the absorption of latent heat in the ethanol,

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which is likely due to a combination of the fact that the latent heat of melting of ethanol is 3X lower than that of water and because the ethanol started at an initial temperature of 70K, far below its melting point of 159K, allowing some heating to occur prior to the transformation.

Fig 4 Temperature comparison between experimental data and model prediction (a) Ice/water case; (b) Ethanol/water case.

Fig 5 Effect of the melting temperature window on the dissolution kinetics Fig. 5 shows a comparison between the results obtained with the model and the experiment for several model runs for the ice/water case. Different runs were made with the model to assess sensitivity to: 1) the temperature range over which the latent heat was absorbed - referred to in the figure as EffCp; and 2) the temperature range over which the material properties of viscosity and permeability (see Eq. 6) where varied - referred to in the figure as MP. The results for varying the EffCp confirm that the discrepancy in plateau temperature is related the phase change temperature range assumed in the model. The more interesting results, however, are those related to varying the MP. As can be seen, the melting process can be delayed by increasing MP. For the case where the MP is increased to 5K. consistent with the latent heat absorption temperature range the melting time is increased beyond 120s. It is believe at this stage that this is related to the formation of a boundary layer at the S/L interface, which when large can significantly impede heat and mass transfer. More work is obviously needed to confirm these results, however, it does make sense from a heat and mass transfer standpoint.

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Conclusion

A coupled thermal-composition-fluid flow model has been developed to simulate the dissolution and homogenization process of a solid alloying element introduced into a liquid at a higher temperature. The impact of thermal and compositionally driven buoyancy and surface tension (Marangoni) flows have been included in addition to a Darcy damping force within the two-phase solid-to-liquid transition regime. This model has been successfully validated against detailed experimental data from an ethanol/ice/water analogue system.

Comparison of data from the ice/water case and ethanol/water cases, show that the combination of thermal and compositional buoyancy and Marangoni forces greatly change the fluid flow pattern and dissolution kinetics in comparison to a system with only thermal buoyancy and Marangoni induced flow. It has also been shown that for the pure ethanol in water case composition buoyancy and Marangoni forces dominate those of thermal buoyancy and Marangoni forces. Work is underway to further validate this model against experimental data obtained with the Ti-Al system.

Reference

[1] R.R. Boyer, An overview on the use of titanium in the aerospace industry, Materials Science and Engineering: A, 213 (1996) 103-114. [2] F.H. Froes, D. Eylon, H.B. Bomberger, Titanium technology: present status and future trends, Titanium Development Association, 1985. [3] A. Mitchell, The electron beam melting and refining of titanium alloys, Materials Science and Engineering A, 263 (1999) 217-223. [4] P.D. Lee, P.N. Quested, M. McLean, Modelling of Marangoni Effects in Electron Beam Melting, Philosophical Transactions: Mathematical, Physical and Engineering Sciences, 356 (1998) 1027-1043. [5] M. Ritchie, S. Cockcroft, A. Mitchell, P. Lee, T. Wang, X-ray-based measurement of composition during electron beam melting of AISI 316 stainless steel: Part I. Experimental setup and processing of spectra, Metallurgical and Materials Transactions A, 34 (2003) 851-861. [6] M. Ritchie, P. Lee, A. Mitchell, S. Cockcroft, T. Wang, X-ray-based measurement of composition during electron beam melting of AISI 316 stainless steel: Part II. Evaporative processes and simulation, Metallurgical and Materials Transactions A, 34 (2003) 863-877. [7] S. Semiatin, V. Ivanchenko, O. Ivasishin, Diffusion models for evaporation losses during electron-beam melting of alphafàeta-titanium alloys, Metallurgical and Materials Transactions B, 35 (2004) 235-245. [8] A.C. Powell, in: Materials Science and Engineering, 1992. [9] K. Vutova, G. Mladenov, Computer simulation of the heat transfer during electron beam melting and refining, Vacuum, 53 (1999) 87-91. [10] S. Akhonin, N. Trigub, V. Zamkov, S. Semiatin, Mathematical modeling of aluminum evaporation during electron-beam cold-hearth melting of Ti-6A1-4V ingots, Metallurgical and Materials Transactions B, 34 (2003) 447-454. [11 ] User Guide for Ansys CFX 12.1, (2009). [12] P. Sahoo, T. Debroy, M. McNallan, Surface tension of binary metal—surface active solute systems under conditions relevant to welding metallurgy, Metallurgical and Materials Transactions B, 19 (1988) 483-491. [13] W. Shyy, Y. Pang, G.B. Hunter, D.Y. Wei, M.H. Chen, Effect of Turbulent Heat Transfer on Continuous Ingot Solidification, Journal of Engineering Materials and Technology, 115 (1993) 8-16. [14] X. Zhao, D.M. Maijer, S.L. Cockcroft, D. Tripp, S. Fox, J. Zhu, in: 11th World Conference on Titanium, Kyoto, Japan, 2007.

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Supplemental Proceedings: Volume 1: Materials Processing and Interfaces TMS (The Minerals, Metals & Materials Society), 2012

COST EFFECTIVE AND ECO-FRIENDLY PROCESS FOR PREPARATION OF WROUGHT PURE Ti MATERIAL

VIA DIRECT CONSOLIDATION OF TiH2 POWDERS

Takanori Mimoto1, Nozomi Nakanishi1, Thotsaphon Threrujirapapong1, Junko Umeda2, Katsuyoshi Kondoh2

'Graduate School of Engineering, Osaka University; 2-1 Yamadaoka; Suita, Osaka 565-0871, Japan

2Joining and Welding Research Institute, Osaka University; 11-1 Mihogaoka; Ibaraki, Osaka 567-0047, Japan

Keywords: Titanium hydride, Pure titanium, Tensile properties, Hot extrusion, Powder metallurgy (P/M)

Abstract

The direct consolidation of T1H2 raw powder in solid-state was employed in order to fabricate powder metallurgy (P/M) pure titanium (Ti) materials. This process consisted of cold compaction, heat treatment and hot extrusion. The raw powder was compacted at room temperature by using conventional press technique at 600 MPa. Their green compacts were dehydrogenated and concurrently sintered into pure Ti billets via the heat treatment process in argon gas atmosphere, and subsequently hot extruded. The microstructural and tensile properties of the wrought extruded specimens were evaluated. The specimen heat treated at 1000 °C for 180 min consisted of the almost same structure, equiaxed et-Ti structure, as a conventional extruded pure Ti, and exhibited the tensile properties of 523 MPa in 0.2%YS, 702 MPa in UTS and 27.1% in elongation at room temperature. These properties satisfied the requirements of Ti Grade 4 of ASTM standard.

Introduction

Titanium materials have recently attracted significant interests as industrial materials because of their excellent characteristics such as low density, high specific strength and superior corrosion resistance. They are widely used in various industrial components, for example, automotive, chemical and marine applications [1, 2]. In particular, Ti materials are very important to reduce the weight of aircraft components from a view point of fuel efficiency improvement [3]. Moreover, they are free from the electrochemical corrosion between carbon composite, such as carbon fiber reinforced plastics (CFRP). Accordingly, Ti materials are suitable for aircraft applications, especially the advanced aircrafts using a large amount of carbon composites, for example, the newest of Airbus A3 80 and Boeing 787. However, the expensive Ti products, resulting from high energy consumption in refining and their poor machinability, prevent their applications from spreading extensively. From a standpoint of these contexts, the cost-effective process to fabricate titanium products is strongly required.

Powder metallurgy process is one of the effective solutions to cut off the high cost of Ti products. P/M process, characterized by near-net-shape capability, provides the three-dimensional complicated component, similar to the end-product after machining, and this leads

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to cost reduction of machining process. The cost reduction effect, compared with the conventional casting process, is over 20% in iron sintered components [4]. Therefore, P/M process is quite suitable for Ti materials with poor machinability from an economical standpoint. On the other hand, the disadvantage in P/M is high-priced starting powder materials. While the total cost of P/M iron sintered components is lower than that of casting products, the cost of starting materials in P/M route is approximately twice as large as that in the casting process is [4]. These facts suggest that reducing raw powder cost gives a great economical advantage to P/M process.

In the present study, extremely low-cost TM2 raw powder, compared with a commercially pure (CP) Ti powder after dehydrogenation treatment, was employed as a promising raw material to fabricate the cost effective P/M pure Ti products. Moreover, the integrated heat treatment process, instead of the conventional dehydrogenation and sintering process, was used in this novel approach in order to reduce the energy consumption and energy cost. Thermogravimetric analysis was carried out on TiH2 raw powder. The effect of temperature at the integrated heat treatment on microstructural and tensile properties of the extruded samples fabricated from TiH2 powder was investigated in detail. The extruded material fabricated from pure Ti powder was employed as a reference specimen. The possibility of cost effective and eco-friendly P/M pure Ti materials via this process was discussed by the above testing and analysis results.

Experimental

TiH2 raw powder (TOHOTEC, TCH450) shown in Fig. 1 (a), having a median particle size of 18.8 urn, was employed as a raw material in this study. The impurity contents of the raw powder were Fe; 0.03, Si; 0.01, Mn; <0.01, Mg; <0.001, Cl; 0.003, N; 0.02, O; 0.13 (in wt.%). The thermal decomposition behavior of the raw powder was investigated by thermal gravimetric analyzer, TG-DTA (SHIMADZU, DTG-60), using heating rate 10 Xï /min from room temperature to 800 t under argon (Ar) gas atmosphere, in order to optimize the heat treatment temperature for dehydrogenation treatment. TiH2 raw powder was compacted at room temperature by applying 600 MPa using a hydraulic direct press machine (Shibayamakikai Co., SHP-200-450) with a press capacity of 2000 kN. These green compacts (H-PC) with a diameter of 41 mm, shown in Fig. 1 (b), were heated in a horizontal tube furnace (Asahi rika, ARF-2-500) at 800 ~ 1000 °C for 180 min in Ar gas atmosphere to remove hydrogen from TiFfe powder and sinter the powder billets, according to the thermogravimetric analysis results. Relative densities of heat treated billets (H-PCHTX: X = 800, 900, 1000) were evaluated by measuring geometric data and mass of the billets. These sintered billets were heated to 1000 °C and kept 3 min under Ar gas atmosphere. The preheated billets were then immediately served to hot extrusion process. The extrusion ratio and speed were 27.9 and 3.0 mm/s, respectively. The mold and die temperature was 400 X^.. The extruded rods had 7 mm in diameter and approximately 600 mm in length. The microstructures and phase characterization of extruded samples (H-PCHTX EXT) were evaluated with optical microscope (OM, OLYMPUS, BX-51P) and X-ray diffraction (XRD, SHIMADZU, XRD-6100). In the sample preparation, the extruded specimens were sectioned parallel to the extrusion direction, abraded with #240 ~ #4000 SiC abrasive papers, and then, buffed with 0.3 urn and 0.05 um AI2O3 particle. After each step, specimen surfaces were carefully rinsed in ethanol with ultrasonic vibration. The extruded rods were machined into tensile specimen bars with 3.0 mm diameter and 20 mm gauge length. The ambient tensile

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properties were evaluated using a universal test machine (SHIMADZU, AUTOGRAPH AG-X 50kN) under a strain rate of 5.0 x 10"4 /s.

As a reference material in evaluating the microstructures and mechanical properties, pure Ti extruded material was also prepared by using pure Ti powder of 99.5% purity (TOHOTEC, TC-450), having a median particle size of 21.9 urn. The powder was consolidated via spark plasma sintering (SPS, SPS Syntex Inc., SPS-1030S) process [5] at 800 *C for 30 min with 30 MPa pressure in vacuum. The sintered pure Ti billet (P-SPS) was extruded in the same way as H-PCHTX EXT specimens. The extruded P-SPS specimen is denoted as P-SPS EXT in this study.

Hydrogen, oxygen, nitrogen and carbon contents of all extruded specimens were measured by hydrogen determinator (LECO Co., RHEN600), inert gas fusion instruments-nitrogen-oxygen determination (LECO Co., TCH600) and carbon-sulfur determination (HORIBA, EMIA-520), respectively.

Figure 1. Morphology of TiH2 raw powder (a) and appearance of TiH2 powder billet compacted at 600 MPa at room temperature (b).

Results and Discussion

Characteristics of TiH? Raw Powder and Its Green Compact

The dehydrogenating temperature of TiH2 raw powder was determined using the result of TG-DTA shown in Fig. 2. The large endothermic peak and weight loss, obviously detected at 400 ~ 700 °C, indicated the thermal decomposition of TiH2 [6-8], Thermodynamic data [9] suggest that TiH2 is completely decomposed into titanium and hydrogen over 773 °C, which is consistent with the result of thermogravimetric analysis shown in Fig. 2. According to the above results, 800 ~ 1000 t were selected for the effective dehydrogenation treatment temperature of H-PCs. In addition, the weight gain, observed over 700 °C, meant an oxidation of pure titanium after dehydrogenation.

Fig. 3 shows the relative density of H-PC, H-PCHT800 ~ 1000, the green compact of pure Ti powder (P-PC) and their heat treated billets (denoted as P-PCHT800 ~ 1000). The P-PC and P-PCHT800 ~ 1000 samples, employing pure Ti powder as a starting material, were prepared via same process as H-PC and H-PCHT800 ~ 1000 as reference specimens. The relative densities of H-PCHT series were significantly improved after heat treatment with higher temperature, although the green density of H-PC was 83.8%, which was lower than that of P-PC 85.7%. In particular, H-PCHT1000 exhibited a relative density of 96.0% and closed-pore

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structures, similar to that of P-PCHT1000 96.7%. Several studies have also mentioned the superior sintering property of T1H2 powder [10-12]. Consequently, the T1H2 powder showed equal or greater sintering property than mat of conventional pure Ti powder. This superior densification behavior of TiH2 compacts, despite the low green density compared to pure Ti powder compact, could be attributed to enhanced mobility of Ti atoms and the contribution which resulted from more vacancies created after hydrogen removal [13]. Moreover, Wang et al. reported the activated surface of Ti powder, transformed from T1H2 powder, was expected to promote its sintering phenomenon [11], and this is certainly related to the deoxidation accelerating the cleaning effect of surface oxide films by liberated atomic hydrogen: Ti02 + 4H -*Ti + 2H20[12].

Figure 2. TG-DTA profiles for TiH2 raw powder as a function of temperature.

Figure 3. Relative density of PCs and PCHTs with various heat treatment temperatures in case of TM2 and pure Ti powders.

Microstructural and Tensile Properties of H-PCHTX EXT Series

Optical microstructures of H-PCHT EXT series (a) ~ (c) and P-SPS EXT (d) are shown in Fig. 4. The area fractions of a-Ti phase (white area) in H-PCHT EXT specimens increased with increase in the heat treatment temperature, and the morphology of a-Ti phase in H-PCHT EXT specimen by applying higher temperature was more similar to that in P-SPS EXT specimen

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in terms of grain shape. In particular, H-PCHT1000 EXT specimen exhibited the almost same structure as P-SPS EXT specimen with equiaxed a-Ti grains. On the other hand, H-PCHT800 EXT specimen obviously consisted of a-Ti phase as a matrix and discrete black area. The grain size of a-Ti phase of this specimen was much finer compared to the other specimens due to the undeveloped a-Ti phase. These fine grains in H-PCHT800 EXT specimen are expected to improve mechanical strength [14-16].

Figure 4. Optical microstructures of H-PCHT EXT series (a ~ c) and P-SPS EXT (d).

X-ray diffraction analysis results for H-PCHT EXT series and P-SPS EXT are presented in Fig. 5. A hydride phase, confirmed with a standard ICDD PDF card, was detected only in H-PCHT800 EXT specimen in addition to the peaks of a-Ti phase. According to the above results, the black area observed in H-PCHT800 EXT specimen shown in Fig. 4 (a) could correspond to hydride phases. It suggests much hydrogen still remained in this specimen because of the inadequate dehydrogenation treatment temperature [17]. The intensity of a peak at 40.45°, indicating a-Ti phase (1 0 -1 1), increased as the heat treatment temperature increased, which is consistent with the OM observations shown in Fig. 4.

Figure 5. XRD profiles of H-PCHT EXT series (a ~ c) and P-SPS EXT (d).

Hydrogen and oxygen content of each specimen shown in Fig. 4 are summarized in Table 1. Hydrogen content of H-PCHT EXT series decreased with increase in the heat treatment temperature. H-PCHT1000 EXT specimen showed the lowest hydrogen content of 0.067 wt.%, which was substantially equal to that of P-SPS EXT specimen with 0.040 wt.%. On the other hand, according to Ti-H phase diagram [18], hydrogen content of H-PCHT800 EXT specimen with 0.33 wt.% means that the growth of a-Ti phase is more difficult compared to the other two specimens using TiH2 powder, corresponding to the microstructures mentioned in Fig. 4. Oxygen

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content of H-PCHT EXT series was higher than that of P-SPS EXT specimen because oxygen content was deeply influenced by a difference of the heat treatment atmosphere between H-PCHT series (in Ar gas atmosphere) and P-SPS (in vacuum atmosphere). In addition, the carbon and nitrogen contents of H-PCHT1000 EXT and P-SPS EXT specimens are as follows; C: 0.007 wt.%, N: 0.024 wt.% (H-PCHT1000 EXT) and C: 0.011 wt.%, N: 0.022 wt.% (P-SPS EXT). Accordingly, with regard to oxygen, carbon and nitrogen contents, H-PCHT 1000 EXT satisfied the requirements of Ti grade 4 of ASTM standard (O: 0.40%max, C: 0.10%max, N: 0.05%max) [1].

Table I. Hydrogen and Oxygen Contents of H-PCHT EXT Series (a ~ c) and P-SPS EXT (d)

Specimen

(a) H-PCHT800 EXT (b) H-PCHT900 EXT (c) H-PCHT 1000 EXT (d) P-SPS EXT

Chemical components (wt.%) Hydrogen

0.33 0.15 0.067 0.040

Oxygen 0.34 0.32 0.32 0.26

Tensile properties of H-PCHT EXT series and P-SPS EXT specimens are summarized in Fig. 6. H-PCHT1000 EXT specimen had 0.2%YS of 523 MPa, UTS of 702 MPa and elongation of 27.1%, which also satisfied the requirements of Ti grade 4 of ASTM standard (minimum YS: 480 MPa, minimum UTS: 550 MPa, minimum elongation: 15%) [1]. However, H-PCHT EXT series exhibited higher strength compared with P-SPS EXT specimen. In particular, H-PCHT800 EXT specimen showed a remarkably high strength of 690 MPa 0.2%YS and 919 MPa UTS, which are 48% and 44% higher than those of P-SPS EXT specimen (0.2%YS: 467 MPa and UTS: 637 MPa), respectively. H-PCHT EXT series were strengthened by solid solution hardening by oxygen [19-21] as mentioned in Table 1. In addition to this effect, the other strengthening effect of H-PCHT800 EXT specimen was probably due to the grain refinement of undeveloped a-Ti phase as mentioned in Fig. 4. It was caused by the uniform distribution of hard hydride phases, as reported by Xu et al. [22].

Figure 6. Tensile properties of H-PCHT EXT series (a - c) and P-SPS EXT (d).

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