Production of high-strength Al-based alloys by

consolidation of amorphous and partially amorphous powders

D I S S E R T A T I O N

zur Erlangung des akademischen Grades

Doctoringenieur (Dr.-Ing.)

vorgelegt

der Fakultät Maschinenwesen

der Technische Universität Dresden

von

Kumar Babu Surreddi

geboren am 1 Febraur 1977 in Visakhapatnam (Indien)

Gutachter: .........................................................

.........................................................

.........................................................

Eingereicht am: .............................................. Tag der Verteidigung: ......................................

II

Table of contents

Acknowledgements ....................................................................................................................IV

Abstract.......................................................................................................................................VI

Introduction ..................................................................................................................................1

Chapter 1: Theoretical background ..............................................................................................5

1.1 Historical background.........................................................................................................5

1.2 Glass formation and crystallization ....................................................................................8

1.2.1 Rapid solidification .....................................................................................................8

1.2.2 Solid state amorphization ..........................................................................................11

1.2.3 Crystallization of metallic glasses .............................................................................14

1.3 Powder metallurgy............................................................................................................21

1.3.1 Powder production.....................................................................................................22

1.3.2 Metal powder compaction .........................................................................................23

1.3.3 Sintering ....................................................................................................................25

Chapter 2: Sample preparation and characterization..................................................................33

2.1 Preparation of amorphous and partially amorphous powders ..........................................33

2.1.1 Melt spinning.............................................................................................................33

2.1.2. Ball milling...............................................................................................................34

2.1.3 Gas atomization .........................................................................................................35

2.2 Powder consolidation .......................................................................................................38

2.2.1 Hot pressing and extrusion ........................................................................................38

2.2.2 Spark plasma sintering ..............................................................................................39

2.3. Sample characterization...................................................................................................40

2.3.1 X-ray diffraction........................................................................................................41

2.3.2 Differential scanning calorimetry..............................................................................41

2.3.3 Parallel plate rheometry.............................................................................................42

2.3.4 Optical microscopy (OM)..........................................................................................43

2.3.5 Scanning electron microscopy (SEM).......................................................................43

2.3.6. Transmission electron microscopy ...........................................................................43

2.4 Mechanical properties ......................................................................................................44

2.4.1 Density.......................................................................................................................44

2.4.2 Microhardness ...........................................................................................................44

III

2.4.3 Compression and tensile testing ................................................................................45

Chapter 3: Characterization of Al-based amorphous powders produced by mechanical alloying

and melt spinning .......................................................................................................................46

3.1 Mechanically alloyed Al85Y8Ni5Co2 glassy powders.......................................................46

3.2 Al85Y8Ni5Co2 glassy powders by milling of melt spun ribbons.......................................51

Chapter 4: Synthesis and characterization of high-strength Al-based alloys by consolidation of

gas-atomized powders ................................................................................................................63

4.1 Gas-atomized Al84Gd6Ni7Co3 powder..............................................................................63

4.2 Gas-atomized Al87Ni8La5 powder ....................................................................................78

4.3 Gas-atomized Al90.4Y4.4Ni4.3Co0.9 powder........................................................................89

Chapter 5: Conclusions and outlook...........................................................................................96

Bibliography…………………………………………………………………………………..100

IV

Acknowledgements

I would first like to express my deep sense of gratitude to my supervisor, Prof. J. Eckert for

providing me an opportunity to be part of his research group at IFW Dresden and for the

continuous support and guidance, without which this work would not have been completed.

His supportive supervision allowed me to work with innovative ideas in a healthy environment.

I am also thankful to him for encouraging me to present my work at several national and

international conferences.

My special thanks go to Dr. Scudino Sergio whose contribution to this thesis is too immense

for words. He is a true researcher with lot of enthusiasm, knowledge, energy and ideas and

inspired me to learn in all aspects. He always willing to help and teach at any time with

continuous support and motivation which allow me to carry my research successfully. Without

his support, this work would not have been possible.

I thank Prof. Dr. Ji-Soon Kim, University of Ulsan, South Korea for allowing me to work with

his group and the kind hospitality during my stay. I thank him for his support to work with

different scientific equipment in his lab mainly with spark plasma sintering.

I would like to gratefully acknowledge Dr. D.J. Sordelet, Caterpiller Inc., Advanced Materials

Technology Group, Mossville, Illinois, USA (previously Ames Laboratory, USA) for

providing me the gas-atomized powders.

I thank Dr. Thomas Gemming and Mohsen Samadi Khoskhoo for their invaluable help with the

transmission electron microscopy investigations.

I thank Birgit Bartusch, Heiko Schulze, Mariana Gründlich, Mihael Frey, and Sven Donath for

their valuable technical help and kindness that made my lab work smooth.

I thank Dr. Horst Wendrock, Dr. Mihai Stoica, Dr. Simon Pauly, Dr. Uta Kühn Dr. Uwe Siegel,

Dr. Vikas Srivastava, Prof. Dr. Gang Liu, Prof. Mariana Calin and Prof. N. K. Mukhopadhyay

for their help and for sharing their knowledge with me.

V

I thank my colleagues and friends who have given an important contribution to this work and

helped me to have conducive environment.

I thank Antje, Christoph, Denise, Dominik, Fahad, Felix, Jayanta, Julia, Kaushik, Marko,

Miroslava, Nilam, Olga, Peter, Prasanth and Shankar for helping me in the work but most

important for the nice and happy time spent together inside and outside the IFW.

The financial support by the Duetscher Akademischer Austausch Dienst (DAAD) for providing

me a scholarship for doing my Ph. D is gratefully acknowledged.

Finally, I am very grateful to my family for their continuous love, encouragement, love and

support.

VI

Abstract

In this thesis, novel bulk Al-based alloys with high content of Al have been produced

by powder metallurgy methods from amorphous and partially amorphous materials. Different

processing routes, i.e. mechanical alloying of elemental powder mixtures, controlled

pulverization of melt-spun glassy ribbons and gas atomization, have been employed for the

production of the Al-based powders. Among the different processing routes, gas atomization is

the best choice for the production of Al-based amorphous and partially amorphous powders as

precursors for the subsequent consolidation step because it allows the production of large

quantities of powders with homogeneous properties (e.g. structure and thermal stability) along

with a uniform size distribution of particles.

Amorphous and nanocrystalline powders have to be consolidated to achieve dense bulk

specimens. However, consolidation of these phases is not an easy task and special care has to

be taken with respect to accurate control of the consolidation parameters in order to achieve

dense bulk specimens without inducing undesirable microstructural transformations (e.g.

crystallization and grain coarsening) or insufficient particle bonding. Consequently, the effect

of temperature on viscosity as well as on phase formation has been studied in detail in order to

select the proper consolidation parameters.

Following their characterization, the Al-based powders have been consolidated into

bulk specimens by hot pressing (HP), hot extrusion and spark plasma sintering (SPS) and their

microstructure and mechanical properties have been extensively investigated. Consolidation

into highly-dense bulk samples cannot be achieved without extended crystallization of the

glassy precursors. Nevertheless, partial or full crystallization during consolidation leads to

remarkable mechanical properties. For example, HP Al84Gd6Ni7Co3 samples display a

remarkably high strength of about 1500 MPa, which is three times larger than the conventional

high-strength Al-based alloys, along with a limited but distinct plastic deformability (3.5 – 4%).

Lower strength (930 MPa) but remarkably larger plastic strain exceeding 25 % has been

achieved for the Al87Ni8La5 gas-atomized powders consolidated by SPS above their

crystallization temperature. Similarly, HP Al90.4Y4.4Ni4.3Co0.9 bulk samples display high

compression strength ranging between 820 and 925 MPa combined with plastic strain in the

range 14 – 30%. Finally, preliminary tensile tests for the Al90.4Y4.3Ni4.4Co0.9 alloy reveal

promising tensile properties comparable to commercial high-strength Al-based alloys. The

mechanical behavior of the consolidated specimens is strictly linked with their microstructure.

VII

High strength and reduced plasticity are observed when a residual amorphous phase is present.

On the other hand, reduced strength but enhanced plastic deformation is a result of the

complete crystallization of the glass and of the formation of a partially or fully interconnected

network of deformable fcc Al.

These results indicate that the combined devitrification and consolidation of glassy

precursors is a particularly suitable method for the production of Al-based materials

characterized by high strength combined with considerable plastic strain. Through this method,

the mechanical properties of the consolidated samples can be varied within a wide range of

strength and ductility depending on the microstructure and the consolidation techniques used.

This might open a new route for the development of innovative high-performance Al-based

materials for transport applications.

1

Introduction

Although annealed commercially pure aluminum is very soft (tensile strength of

about 50 MPa) [Rooy 1990], the strength of Al can be increased by adding proper alloying

elements and by exploring versatile strengthening methods such as solid solution

hardening, work hardening and precipitation hardening according to the specific

application and usage [Cahn 1993]. For example, the high-strength Al alloy ‘7075-T6’

(with Zn, Mg and Cu as major alloying elements and with specific temper treatment ‘T6’)

reaches a value of tensile strength of 540 MPa [Kaufman 2000]. Due to their remarkable

properties, such as high specific strength, good corrosion resistance and electrical

conductivity, along with excellent machinability and formability, Al-based alloys have

become very popular in many engineering applications [Rooy 1990]. In particular, high-

strength Al-based alloys are gradually substituting steel and cast iron in automotive and

aerospace industries to reduce the structural weight of vehicles in order to decrease fuel

consumption and, consequently, environmental pollution [Davis 1993]. Besides high-

strength Al-based alloys strengthened by conventional methods, a new class of lightweight

high-performance Al-based alloys, the Al-based amorphous and nanocrystalline materials,

have been recently developed.

The starting point for the development of amorphous, partially amorphous and

nanocrystalline Al-based alloys as potential candidates for structural as well as functional

applications was the discovery of high tensile strengths, σf, exceeding 1200 MPa for melt-

spun Al-RE-TM (RE = rare earth, TM = transition metal) amorphous ribbons or

amorphous wires prepared by melt-extraction [He 1988, Inoue 1994a, Inoue 1988c, Tsai

1988]. Such strength levels are about twice as high as for conventional crystalline Al alloys

[Inoue 1998]. Even higher strength levels are observed for nanostructured two-phase alloys

with nanoscale fcc Al particles in an amorphous matrix. Such nanoscale mixed-phase

materials can be made directly upon rapid quenching or by partial devitrification of the

melt-spun amorphous ribbon. They exhibit tensile fracture strengths reaching up to

1560 MPa for an Al88Ni9Ce2Fe1 alloy [Kim 1991], which is considerably larger than that of

the corresponding single-phase amorphous alloy (1100 MPa). The tensile strength

increases almost linearly with the volume fraction of particles, Vf, reaching its maximum

value for Vf = 25%. Simultaneously, the Vickers hardness increases from 280 to 400, and

the Young’s modulus increases from 63 to 71 GPa with increasing Vf [Kim 1991].

2

Although Al-based amorphous and nanocrystalline alloys exhibit improved

mechanical properties compared to conventional Al-based crystalline alloys, the maximum

scale of the products is limited to a thickness of less than 100 micrometers due to their

relatively low glass-forming ability [Inoue 1998]. In general, Al-based metallic glasses can

only be obtained by melt-spinning in the shape of ribbons or by gas atomization in form of

powder [Inoue 1998]. Only very recently, a single-phase amorphous Al-based alloy with

1 mm diameter has been successfully produced by solidification methods for the

composition Al86Ni8Y6 [Yang 2009]. This limitation has prevented a wide extension of

application fields of the Al-based amorphous alloys even despite their excellent

mechanical properties [Inoue 1998].

To overcome this limitation, powder metallurgical methods, such as gas

atomization followed by hot extrusion, have been employed to create bulk Al-based

samples with the desired microstructure [Inoue 2001a, Kawamura 2001, Ohtera 1991,

Ohtera 1992]. Amorphous alloys with Al concentrations above about 85 at.% exhibit a

glass transition phenomenon, followed by a rather narrow supercooled liquid region of

about 20 K prior to crystallization [Inoue 1988c]. This promises that by utilizing the

viscous flow of the supercooled liquid, bulk amorphous samples can be produced by

consolidation of mechanically alloyed or gas-atomized powders [Inoue 1988a]. As an

example, Al85Ni10Y5 amorphous alloy powders were produced by high pressure helium

atomization [Afonso 2003, Inoue 1988c]. The DSC trace from the Al85Ni10Y5 amorphous

alloy powder exhibited the same thermal stability as that for the as-quenched ribbons, and

no appreciable difference in Tg [Inoue 1988a]. Bulk Al85Ni10Y5 amorphous alloy powders

through warm extrusion at extrusion ratios 4 and 7 were tried by changing the extrusion

temperature [Kato 1994]. By this, full density of Al85Ni10Y5 amorphous alloy powders was

obtained [Kawamura 1993]. Recently, the development of amorphous Al85Y8Ni5Co2 alloy

powders by mechanical alloying was carried out [Börner 2001]. The amorphous

Al85Y8Ni5Co2 alloy powders exhibited good thermal stability and low viscosity in the

supercooled liquid region. Vickers hardness (HV) values of about 430 were found for

mechanically alloyed and consolidated Al85Y8Ni5Co2 [Börner 2001]. Neglecting residual

porosity in the bulk samples, the fracture strength is estimated to be about 1400 MPa.

However, there were still problems to consolidate the powders to high density.

Alternatively to mechanical alloying of elemental powder mixtures, Al-based

glassy powders can be produced by controlled milling of melt-spun glassy ribbons. For

example, Al85Ni9Nd4Co2 with a thickness of ≈ 25 µm and a width of ≈ 6 mm were

3

planetary ball-milled at cryogenic temperatures [Calin 2004]. By this glassy Al-powder can

be produced in bigger quantities and the sticking problems, which normally occur when

one tries to mechanically alloy the pure elements with this composition (due to the high Al

content of 85 at.%) can be avoided. Subsequent uniaxial hot-pressing leads finally to a

density of about 94% in the bulk sample without crystallization [Calin 2004].

Powders can be consolidated by a variety of well−known and developed techniques,

including cold−pressing followed by high−temperature sintering, cold and hot extrusion

and hot isostatic pressing [Koch 2002]. Less common consolidation techniques are spark

plasma sintering [Mamedov 2002] or microwave [Yadoji 2003] and laser [Singh 2009]

assisted sintering, which involve the use of pulsed electric current, microwave and laser

irradiation to improve powder consolidation. These techniques are carried out at relatively

lower temperatures for a shorter time than in conventional sintering processes. Therefore,

they show a large potential for achieving fast and full densification of metastable materials

(e.g. amorphous or nanostructured phases). Regardless of the processing route used, the

essence of any compaction technique is to apply high pressure and high temperature to

achieve full density with minimal grain growth and/or undesirable microstructural

transformations [Koch 2002]. The choice of the consolidation method is usually dictated

by the morphology of the powders, their intrinsic characteristics (yield strength, chemical

diffusivity) and the need to prevent or induce structural changes in the compacted product

brought about by grain growth or phase transformations [Koch 2002]. Consolidation of

metastable phases is not a trivial process and often results in phase transformations (e.g.

crystallization and grain coarsening), or insufficient particle bonding. For example,

amorphous alloys, being thermodynamically metastable with respect to the crystalline state,

cannot be heated at elevated temperatures for long times without running the risk of

inducing crystallization. Similar difficulties arise in the consolidation of nanocrystalline

materials, where the diffusional processes will not only assist densification, but will also

allow grain coarsening to occur [Koch 2002]. These characteristics severely limit the

consolidation parameters that can be used and, as a result, temperature, pressure and the

time span of the consolidation process have to be adjusted carefully in order to achieve a

balance between good densification and desired microstructure.

The aim of this thesis is to develop novel high-strength nanocrystalline Al-based

alloys with high Al content. In order to achieve this purpose, highly dense bulk

nanocrystalline samples have been prepared by powder metallurgy from Al-based

amorphous and partially amorphous precursors. The work is focused on three specific

4

aspects: (i) production and characterization of Al-based amorphous and partially

amorphous powders by different processing routes, (ii) consolidation of the powder

precursors into bulk samples with the desired microstructure and (iii) evaluation of the

mechanical properties of the consolidated specimens.

A theoretical background is given in Chapter 1 along with an overview of the basic

concepts of rapid solidification, solid state amorphization and crystallization behavior of

amorphous metallic alloys. In addition, the main aspects of powder production, powder

compaction and sintering are also presented in this chapter. A summary of the

experimental procedures used in the course of this work, which includes details on the

processing routes and characterization techniques, is given in Chapter 2. The effect of the

processing route on thermal stability as well as on the structure and crystallization of the

materials prepared by mechanical alloying and mechanical milling of melt-spun ribbons

are discussed in Chapter 3 along with the consolidation of the bulk samples and the

corresponding mechanical behavior. In Chapter 4, a detailed characterization of the

structure, thermal behavior and mechanical properties of gas-atomized powders is

discussed. Finally, summary and conclusions are presented in Chapter 5.

Chapter 1: Theoretical background

In this chapter the necessary information related to this work is presented. The

initial part of this chapter deals with the formation of amorphous metallic alloys produced

by different techniques and includes a brief description of the historical development of

metallic glasses. Details about nucleation and growth mechanisms, which are very

important aspects to understand the formation and the crystallization of metallic glasses,

are also given. In the second part of this chapter, basic concepts about powder metallurgy

related to production and consolidation of amorphous powders is presented along with the

background information about development and properties of Al-based amorphous alloys

and composites.

1.1 Historical background

In 1960 Klement, Willens, and Duwez [Duwez 1960] first successfully applied a

rapid cooling method, known as splat quenching, to form an Au-Si amorphous alloy. Very

high cooling rates of the order of 107 K/s were used to cool the Au-Si liquid alloy from

1300°C to room temperature. Soon after, Chen and Turnbull [Cohen 1961] suggested that

the formation of an amorphous phase in Au75Si25 was connected with the existence of a

deep eutectic near this composition. The presence of a deep eutectic gives the melt the

opportunity to cool to a temperature at which its viscosity is quite high and, consequently,

the atomic mobility in the melt has been reduced, thus hindering nucleation. This idea led

to the discovery of glass formation in a number of alloys, including Pd-Si [Chen 1969], Ni-

Nb [Ruhl 1967], Fe-C-P [Duwez 1967], Zr-TM (TM=Ni, Pd, Cu, Co) [Ray 1968] Pd-Si

[Chen 1969]. Later, it was found that the thickness of the glassy materials can be

remarkably increased by the addition of solute elements to the binary alloy systems. For

example, Chen [Chen 1974] successfully produced Pd-M-P (M = Ni, Co and Fe) glassy

alloys with diameter of 1-3 mm. The critical diameter for glass formation was

progressively increased in the following years. In the early 1980’s Turnbull’s group was

able to produce glassy ingots of Pd40Ni40P20 with diameter of 5 mm [Drehman 1982]. In

1984, the critical diameter was further extended to 10 mm by using boron oxide flux to

purify the melt and to eliminate nucleation [Kui 1984].

6

A great impulse in the development of bulk metallic glasses (BMGs) was given by

Inoue and coworkers. In the late 1980s, they observed a large super-cooled liquid (SCL)

region of about 70 K in Al-La-Ni alloys [Inoue 1989b]. This wide interval implies that the

SCL region can exist in a large temperature range without crystallization. They proposed

that such a high resistance to crystallization at temperatures above the glass transition

might reflect a lower critical cooling rate and, consequently, a high glass-forming ability

(GFA). Therefore, assuming a link between high ΔTx values and high GFA, they looked for

alloys with large ΔTx. Indeed, they found a number of Mg- and Zr-based alloys with large

values of ΔTx, which can be cast into fully glassy rods with a thickness of several

millimeters [Inoue 1992, Inoue 1993, Zhang 1991]. Over the last four decades, the critical

casting thickness has been increased by more than three orders of magnitudes and

amorphous components with mass of several kilograms can be fabricated. To date, more

than a thousand different bulk metallic glasses have been produced in Zr-, Fe-, Pd-,Ni-,

Cu-, Mg-, and Ti-based systems. In 1993, Johnson and Peker developed a multi-component

Zr41.2Cu12.5Ni10Ti13.8Be22.5 metallic glass with a critical cooling rate of 1 K/s [Peker 1993].

This alloy became the first commercial bulk metallic glass and is known as Vitreloy 1.

The first experiments on the formation of Al-based amorphous alloys were done by

liquid quenching for binary Al-metalloid and Al-TM (TM = transition metal) alloys.

Coexistence of amorphous and crystalline phases was found in Al-Si [Predecki 1965], Al-

Ge [Laridjan 1972] and Al-TM (TM = Cu [Davies 1972], Cr [Furrer 1977] or Pd [Sastry

1981] alloys). In 1981, a single phase amorphous alloy with more than 50 at.% aluminum

was produced in the ternary alloy systems Al-Fe-B and Al-Fe-Ge [Inoue 1981]. However,

these alloys showed extreme brittleness. Similar results were obtained for Al-Fe-(Si/Ge)

and Al-Mn-Si alloys [Inoue 1987, Suzuki 1983] and it was proposed that the brittleness

might be an inherent property of Al-based amorphous alloys. The concept of inherent

brittleness of Al-based amorphous alloys was changed with the discovery of an Al-Ni-

(Si/Ge) amorphous alloy with more than 80 at.% Al characterized by good bending

ductility [Inoue 1987]. Since then, ductile Al-based amorphous alloys have been found in

several binary and ternary alloy systems, such as Al-RE binary alloys (RE = rare earth)

[Inoue 1988b, Inoue 1989a], Al-ETM-LTM (ETM = early transition metal; LTM = later

transition metal) [Inoue 1988b, Tsai 1988] and Al-RE-LTM ternary alloys [Inoue 1988b,

Tsai 1988].

7

Among the different Al-based glassy alloys, Al-Y-Ni amorphous alloys are of

particular interest due to their interesting mechanical properties. For example, Inoue et al.

[Inoue 1988b] and Poon and coworkers [He 1988] reported that Al-Y-Ni amorphous alloys

produced by melt spinning display high strength together with good bending ductility. The

amorphous phase was formed over a wide composition range encompassing 3–22 at.% Y

and 4–33 at.% Ni. The addition of Co to the Al-Y-Ni glass forming system leads to

increased mechanical strength with no detriment to the bending ductility [Inoue 1990]. For

example, the Al85Ni5Y8Co2 (at.%) glassy alloy displays high tensile fracture strength of

1250 MPa.

Even higher strength levels can be achieved through the development of a

nanostructured two-phase microstructure consisting of a homogeneous dispersion of

nanoscale fcc Al particles in the amorphous matrix. In general, Al–rich amorphous alloys

(Al > 85 at.%) crystallize through two exothermic reactions. The first reaction is due to the

precipitation of nanocrystalline fcc Al particles and the second results from the

decomposition of the remaining amorphous phase to intermetallic compounds [Inoue 1998].

Therefore, by annealing the glass at temperatures above the first crystallization event, a

two-phase microstructure containing nanoscaled fcc Al particles can be formed. The

annealed alloy exhibits high tensile strength exceeding 1400 MPa which is about 1.4 times

higher than the corresponding single-phase amorphous alloy [Inoue 1994b]. A similar

microstructure can also be formed by varying the cooling rate during glass formation.

When the cooling rate of the Al-rich alloys is controlled, it is possible to obtain a nanoscale

mixed structure consisting of fcc Al particles with a size of 3 – 5 nm embedded in an

amorphous matrix of the melt spun ribbons [Kim 1991]. The volume fraction of the fcc Al

phase is controlled by changing the rotation speed of the melt spinning wheel [Inoue 1998].

The increase in strength by the homogeneous dispersion of nanoscale fcc Al

particles in the residual amorphous matrix can be related to the following factors [Inoue

1997a]: (1) defect-free nanoscale fcc Al particles (Al particles are too small to contain

dislocations), (2) interface effect, due to the highly dense packed atomic configuration of

the amorphous/Al particle interface, and (3) nanoscale effect, related to the particle size of

the Al particle that is smaller than the width of the inhomogeneous shear deformation

region and, as a result, the nanoscale Al particles can act as an effective barrier against the

shear deformation of the amorphous matrix [Inoue 1998].

8

1.2 Glass formation and crystallization

Generally, a glassy or amorphous phase can be formed through two ways, i.e.,

through liquid-to-solid and through solid-to-solid transformation. The liquid-to-solid

transformation, e.g., rapid solidification, consists of cooling a liquid so quickly that

crystallization has not enough time to occur. On continued cooling, as the temperature is

lowered, the viscosity of the liquid increases and the atomic mobility decreases. At low

cooling rate the liquid is able to crystallize, but crystallization can be avoided if the cooling

is sufficiently rapid, resulting in a progressive freezing of the liquid configuration [Scherer

1993 ]. That is, the viscosity becomes so high and the atomic mobility so slow that the

liquid cannot change its structure rapidly enough to stay in internal equilibrium and,

ultimately, the final product is a solid with a liquid-like atomic arrangement [Cahn 1996].

In contrast, the solid-to-solid transformation, e.g., solid-state amorphization by mechanical

alloying of elemental powder mixtures [Hellstern 1986, Koch 1983, Schwarz 1985],

involves an increase in the energy of the starting crystalline material by the addition of

some externally provided energy, and the storage of this energy in the crystal up to a point

at which it becomes unstable with respect to the amorphous state. The highly energized

material then lowers its energy by transforming into a different atomic structural

arrangement, i.e., the glass [Schultz 1994].

1.2.1 Rapid solidification

Rapid solidification (RS), which includes techniques such as melt spinning and gas

atomization, involves high velocity of propagation of the advanced solidification front

[Jones 1999] and cooling rates from 103 to 109 K/s [Anantharaman 1987]. Rapid

solidification results from the rapid extraction of the heat of transformation from the mass

of molten metal or alloy either directly by the external heat sink and /or internally by the

undercooled melt. When extraction of heat is so rapid, the liquid undergoes a significant

undercooling [Anantharaman 1987] and, as a result of the limited atomic mobility, the

structural disorder of the liquid phase liquid is retained (quenching-in) in the glassy

material. Therefore, as a first approximation, the structure of a glass can be considered as a

frozen liquid.

The main features of this process can be understood by considering the changes in

Gibbs free energy G, viscosity η and density ρ, which occur when a glass-forming system

9

is cooled from the liquid into the crystalline or glassy states (Figure 1.1 free energy,

viscosity and density diagram). At high temperatures, above Tliq, the liquid has a lower G

than the crystalline solid. Therefore, it is the thermodynamically stable phase at that

temperature. Below Tliq the undercooled melt is thermodynamically less stable than the

crystalline phase and may crystallize if a critical nucleus is provided, giving rise to the

discontinuous increase of viscosity shown in Figure 1.1.

Figure 1.1 Schematic representation of the variation in Gibbs free energy G, viscosity η and density ρ as a function of temperature occurring when a glass-forming system is cooled from the liquid state into the crystalline or glassy states (after [Cahn 1996, Davies 1983, Greer 1993]).

On the other hand, crystallization can be avoided if the cooling rate is high enough

to prevent nucleation. In this case, the system continues to follow the liquid Gibbs free

energy curve without any change at Tliq. The viscosity continuously increases and its

10

equilibrium values can be well described by the Vogel-Fulcher-Tamman (VFT) empirical

equation [Fulcher 1925, Tammann 1926, Vogel 1921] :

⎟⎟⎠

⎞⎜⎜⎝

⎛−

=0

0 expTT

bηη . (1.1)

In this region, the viscosity shows a strong temperature dependence, reflecting the

atomic configurational changes in the liquid. As a response to the imposed cooling rate, the

system progressively relaxes into equilibrium configurations by modifying the relative

atomic positions in order to minimize the Gibbs free energy at each temperature. These

modifications have consequences on the physical properties of the supercooled liquid, such

as density and viscosity, which increase with decreasing temperature [Cahn 1996].

The atomic mobility decreases markedly as density and viscosity increase. The

reduction of the atomic mobility due to the rapid increase of the liquid viscosity

progressively extends the average time required for the rearrangement of the atomic

positions into the equilibrium configuration (the structural relaxation time). Finally, at the

so-called glass transition temperature Tg, the atomic mobility is so slow that there is no

time for configurational changes, and below Tg the system is frozen into a non-equilibrium

configuration, which corresponds to the equilibrium configuration at Tg [Scherer 1993].

The lack of configurational changes of the glass is revealed by the weak temperature

dependence of the viscosity below Tg, which departs from the equilibrium values predicted

by the VFT equation and is similar to that characterizing the crystal.

Since the transformation from the liquid to the glassy state is progressive and

continuous, it might be expected that the structure of the two states would be similar. This

is confirmed by diffraction experiments [Egami 1993, Finney 1977]. The diffraction

intensities of amorphous alloys are rather similar to those of the liquid, namely diffuse

diffraction peaks without discrete lines, which characterize a crystalline substance, thus

indicating the similarity of the atomic structure of these two states [Egami 1993]. Like

liquids, glasses possess a disordered structure lacking long-range order. That is, in a glass,

there is no regular arrangement resulting from the distribution over long distances of a

repeating atomic arrangement, characteristic of a crystal. There is only evidence of a short-

range order, which corresponds to the mutual arrangement of the nearest neighbors to a

given atom and varies according to the atomic site considered [Elliott 1990].

11

1.2.2 Solid state amorphization

Alternative to rapid solidification methods, amorphization can be achieved in the

solid state by ball milling (BM), a group of techniques that combines deformation,

comminution and mixing [Cahn 1993]. Solid-state amorphization by ball milling is mainly

divided into two different routes depending on the starting material. Ball milling of

powders with different compositions (mixture of elemental powders as well as of

intermetallic compounds), in which material transfer occurs, is named mechanical alloying

(MA), while ball milling of single composition powders, such as single-phase compounds,

where material transfer is not required, has been termed mechanical milling (MM) [Koch

1996].

Mechanical alloying: The basic principles of solid-state amorphization by mechanical

alloying are described in Figure 1.2. The thermodynamic stable state of a system is

determined by the minimum in the Gibbs free energy G. Below the melting temperature,

the equilibrium crystalline state is characterized by a Gibbs free energy Gx, which is

always lower than that of the glass. The amorphous state is a metastable state, i.e. an

energy barrier prevents the amorphous phase from spontaneous crystallization.

Figure 1.2 Schematic representation of the basic principles of amorphization by solid-state reaction (after [Schultz 1988, 1994]).

In order to achieve the amorphous phase by solid-state reaction, it is necessary first

to establish an initial state of high free energy G0. For the mechanical alloying process, this

state consists of the layered system of two crystalline system of two elemental metals

[Schultz 1994]. The free energy of the system can be then lowered from G0 either by the

12

formation of the metastable amorphous phase with free energy Gam or by the formation of

the equilibrium crystalline phase. The crystalline phase is thermodynamically favored

since the driving force for crystallization ΔGx = (G0 - Gx) is larger than that for

amorphization ΔGam = (G0 - Gam). However, the formation of crystalline or amorphous

phases depends on thermodynamic as well as on kinetics factors. The formation of the

amorphous phase is then possible if the system is kinetically restricted from reaching

crystalline states of lower free energy, i.e. the amorphization reaction proceeds

considerably faster than the formation of the equilibrium crystalline phase from the initial

state G0. At the same time, the metastable amorphous phase must not crystallize as the

reaction proceeds, i.e. the timescale for crystallization of the amorphous phase must be

longer than that of amorphous phase formation. These kinetic constraints can be

summarized as [Schultz 1988, 1994]:

am→0τ << x→0τ , (1.2)

am→0τ << xam→τ , (1.3)

where am→0τ and x→0τ are the characteristic reaction times for the formation of the

amorphous and equilibrium crystalline phase, respectively, and xam→τ is the timescale for

crystallization of the amorphous phase.

The process of mechanical alloying is schematically shown in Figure 1.3. During

milling the balls collide with each other and with the internal side of the milling container

and a certain amount of powder is trapped between colliding balls during each impact. Due

to the heavy deformation, in the early stages of milling the particles are cold-welded and

plastically deformed, leading to the characteristic layered structure consisting of various

combinations of the starting constituents [Schultz 1994, Suryanarayana 2001], as

illustrated in Figure 1.3. With increasing milling time the thickness of the individual layers

decreases. The structure undergoes severe deformation, work hardening and fracture.

Fragments generated by this mechanism may continue to reduce in size, giving rise to a

more and more refined microstructure. However, the structure cannot be refined

indefinitely by deformation due to the increasing hardness with decreasing crystallite size,

meaning that ball milling does not lead to mixing at an atomic scale, but produces an

ultrafine layered microstructure [Hellstern 1989, Schultz 1994].

13

 

Figure 1.3 Ball-powder-ball collision of powder mixture and formation of layered powder particles during mechanical alloying ([Suryanarayana 2001]).

Alloying begins to occur at this stage due to the combination of decreased diffusion

distances (interlayer spacing), increased lattice defect density, and any heating that may

have occurred during the milling operation [Suryanarayana 2001]. Therefore,

amorphization during MA is not a purely mechanical process, but involves an

interdiffusion process, driven by the negative heat of mixing, of the thin layers in a similar

way as observed in diffusion couples [Schwarz 1983].

Mechanical milling: Amorphization by MM consists of energizing the equilibrium

crystalline compound by the severe cyclic deformation provided by the milling process.

The advantage of MM over MA is that since the powders are already alloyed and only a

reduction in particle size and/or other transformations need to be induced mechanically, the

time required for processing is generally shorter than that for MA [Suryanarayana 2001].

An interesting aspect of MM is that instead of lowering the Gibbs free energy of the

system, in this process the free energy of the equilibrium crystalline compound is raised to

a level equal to or larger than that of the amorphous phase. The mechanical treatment

increases the Gibbs free energy of the intermetallic compound by the generation of

chemical disorder, point defects, such as vacancies, and lattice defects (e.g. dislocations)

[Schwarz 1988]. In addition, an important contribution to the energy increase most likely

comes from the reduction of grain size to a nanometer level and the consequent storage of

energy in the grain boundaries, which constitute an appreciable fraction of the material

volume [Bakker 1995].

14

1.2.3 Crystallization of metallic glasses

Regardless of the processing route used for their production, metallic glasses are

not in a state of internal equilibrium and, when heated to a sufficiently high temperature,

they tend to a more stable condition. Upon annealing below the glass transition

temperature, the glass initially relaxes towards a state corresponding to the ideal frozen

liquid with lower energy [Cahn 1993]. The structure evolves to one with higher density,

which could be considered characteristic of glass formation at a slower cooling rate [Cahn

1993] and finally, above the glass transition temperature, the glass crystallizes.

Crystallization studies of metallic glasses are of primary importance not only in

order to analyze their thermal stability against crystallization but also to investigate the

fundamental aspect of the processes of nucleation and growth. Metallic glasses crystallize

by a nucleation and growth mechanism [Köster 1981] in a similar way as for solidification

of liquids below their liquidus temperature; however, since the crystallization process is

much slower than solidification of liquids, it is relatively easier to investigate the

crystallization in glasses than in liquids. Metallic glasses can be considered as deeply

undercooled liquids, therefore, their crystallization behavior may be analyzed in a similar

way.

When a liquid is cooled below the liquidus temperature, it is energetically less

stable than the crystalline phase and tends to transform to the more stable crystalline solid

(Figure 1.4). The difference in Gibbs free energy between these phases provides the

driving force for the nucleation process. However, the crystallization does not start

immediately after the system has reached the range of parameters where the new phase has

the lowest free energy. The liquid must be undercooled below Tliq before crystallization

can occur, due to the existence of an energy barrier to nucleation [Herlach 1997, Kelton

2004]. The crystallization of a liquid is not a transformation that occurs in the entire

volume at once, but it starts and progressively extends from discrete centers throughout the

material [Christian 2002a]. These centers are aggregates of atoms characterized by an

atomic configuration similar to the lattice of the product phase. However, not all these

aggregate s of the new phase, called embryos, are stable. In fact, embryos below a critical

15

Figure 1.4 Schematic free energy diagram as a function of temperature for a liquid undercooled below the liquidus temperature (Tliq). Gliq and Gsol refer to the Gibbs free energies of the liquid and solid phases, respectively (after [Porter 2009]).

minimum size are associated with an increase in Gibbs free energy, are instable, and thus

quickly disintegrate [Burke 1965]. The reason for this is that the formation of an embryo

within the parent phase is accompanied by the creation of an interface [Burke 1965,

Christian 2002a]. Due to the different structures between the liquid and the crystal, a

mismatch along the interface arises. The positive energy associated to this interface has to

be supplied by the Gibbs free energy of the transformation and thus disfavors the

crystallization of the liquid [Burke 1965]. On the other hand, aggregates larger than the

critical size are stable and capable of continuous existence. Such stable structures are

called nuclei and their formation is termed nucleation [Porter 2009]. Nucleation that occurs

randomly throughout a system in the absence of foreign bodies that can catalyze

crystallization is said to be homogeneous. In contrast, nucleation at preferred sites is named

heterogeneous [Turnbull 1950].

The resistance of liquids to nucleation can be better understood in the framework of

the classical theory of nucleation [Christian 2002a] for vapor condensation, where it is

assumed that the embryos have uniform structure, composition and properties. These

assumptions leave the shape and size of the embryo or nucleus as the only variable

parameters. For homogeneous nucleation, ΔG of formation of a spherical embryo of radius

r within the liquid phase is given by [Burke 1965, Christian 2002a, Fisher 1948, Porter

2009, Turnbull 1969]:

16

ΔG = 3

4π r3 ΔGv + 4π r2 σ , (1.4)

where ΔGv is the Gibbs free energy difference between unit volume of the crystal and

liquid, and σ is the interfacial energy per unit area of the solid/liquid interface. Figure 1.5

shows ΔG as a function of r at a temperature below Tliq, where the crystal is

thermodynamically more stable than the liquid, i.e. ΔGv is negative below Tliq (ΔGv and σ

are assumed to be independent of r). There is a clear competition between the interfacial

energy and the ΔGv terms. In fact, the interfacial energy is always positive and, therefore,

is opposed to ΔGv. ΔG passes through a maximum, denoted ΔG#, at a radius r# (the critical

nucleus size). In the case of embryos with radii smaller than r#, the interfacial energy is

greater than the volume free energy, with the result that there is a net increase in ΔG upon

growth and the embryos have the tendency to shrink rather than to grow. In contrast, for

values greater than r#, the volume free energy term dominates on the surface term because

it is proportional to r3. In this case, the net free energy change accompanying the

transformation is negative, with the result that large embryos (nuclei) are stable. Embryos

of radius r# have an equal possibility to shrink or to grow [Burke 1965, Porter 2009].

Figure 1.5 Free energy of formation of a spherical embryo as a function of the radius r (after [Herlach 1993]).

The critical nucleus size, r#, is defined by the condition 0#

=∂Δ∂

=rrrG , which gives

17

r# = GΔ

−σ2 . (1.5)

The critical value of ΔG corresponding to r# is equal to

ΔG# = 2

3

316

vGΔπσ , (1.6)

which corresponds to the activation energy for homogeneous nucleation, i.e. the barrier to

nucleation that has to be overcome in order to form a nucleus of critical size [Burke 1965,

Porter 2009].

Metallic glasses can be used as precursors for nanocrystalline materials, perhaps the

most attractive microstructure from the point of view of the functional properties. For

example, glass-matrix composites consisting of nanosized particles embedded in a glassy

matrix can be produced by controlled devitrification (crystallization) of metallic glasses

[Chen 1999, Inoue 1997b, Inoue 1999]. This technique has been used for long time for

conventional glasses [Holand 2001] in order to produce composite materials with a wide

variety of microstructures and advantageous properties.

The basic principle for the production of glass-matrix composites by crystallization

of a glassy precursor is to control the crystallization kinetics by optimizing the annealing

conditions (annealing temperature and time, heating rate, etc.) and chemical composition in

order to obtain a glassy phase that partially or completely transforms into a nanocrystalline

material with the desired microstructure [Cahn 1996, Köster 1981].

Controlling the microstructure development from amorphous precursors requires

detailed understanding of the specific mechanisms influencing structural transformations.

Thermal analysis, in particular differential scanning calorimetry (DSC), has been

successfully employed for studying the phase transformations involving nucleation and

growth, continuous grain growth of pre-existing nuclei and in general for investigating the

crystallization kinetics of glass-forming liquids and metallic glasses [Scott 1977, Weinberg

1996, Yinnon 1983].

Crystallization is a thermally activated reaction. A general objective of the

modeling of thermally activated reactions is the derivation of a complete description of the

progress of a reaction that is valid for any thermal treatment, be it isothermal or by linear

heating (isochronal). However, this is a difficult task because any reaction might progress

through a range of mechanisms and intermediate stages, all of which can have a different

18

temperature dependency. To come to terms with this potentially very complicated problem,

most researchers attempt to achieve this objective by making a few judiciously simplifying

assumptions. A simplifying assumption that is encountered in numerous publications is the

hypothesis that the transformation rate during a reaction is the product of two functions,

one depending solely on the temperature, T, and the other depending exclusively on the

fraction transformed, α [Starink 2004].

( ) )(Tkfdtd αα

= . (1.7)

The temperature dependent function is generally assumed to follow an Arrhenius-

type dependency

⎟⎠⎞

⎜⎝⎛−=

RTEkk exp0 . (1.8)

Thus, to describe the progress of the reaction at all temperatures and for all

temperature-time programs, the function f(α), the reaction constant, k0 and the activation

energy, E need to be determined. In general, the reaction function f(α) is unknown at the

outset of the analysis. From the above equations, it follows that for transformation studies

performed at a constant temperature, T, E can be obtained from the equation [Christian

2002b]:

( ) ii

f cRTEt +⎟⎟

⎠

⎞⎜⎜⎝

⎛=ln , (1.9)

where tf is the time needed to reach a certain fraction transformed and ci is a constant,

which depends on the reaction stage and on the kinetic model. Thus, E can be obtained

from two or more experiments at different T. For isothermal experiments, k(T) is constant,

the determination of f(α) is relatively straightforward, and is independent of E.

For non-isothermal experiments, the reaction rate at all times depends on both f(α)

and k(T), and the determination of f(α), k0 and E (the so-called kinetic triplet) is an

interlinked problem. A deviation in the determination of any of the three will cause a

deviation in the other parameters of the triplet. Over the past decades a variety of non-

isothermal methods have been proposed. Though most of them are used for oxide-glasses,

in case of metallic glasses the most widely used non-isothermal methods are the Kissinger

19

analysis [Kissinger 1957], Ozawa analysis [Ozawa 1965], Matusita and Sakka analysis

[Matusita 1984], Gao and Wang method [Gao 1986], Augis and Bennet’s method [Augis

1978], and Lasocka’s method [Lasocka 1984]. While isothermal analyses are in most

cases more definitive, it has been shown that the non-isothermal technique also has several

advantages, in particular that experiments can be performed quite rapidly [Henderson

1979]. Additionally, many phase transformations occur too rapidly to be measured under

isothermal conditions because of transients associated with the experimental apparatus

[Henderson 1979]. In this thesis, the isochronal method employed for calculating the

activation energy for the crystallization has been the Kissinger method. Although, the

Kissinger analysis was not originally developed for solid-state transformations, Henderson

has shown that it is applicable to these transformations [Henderson 1979]. The activation

energy calculated using the Kissinger approach depends on the temperature dependences

of the nucleation and growth rates and on any transient events, which they may exhibit.

Despite difficulty in interpretation, this approach has been widely used for comparing the

stability of metallic glasses [Louzguine 2002a, Stoica 2009, Venkataraman 2007a,

Venkataraman 2005]. A high value of the activation energy is generally interpreted as a

measure of the high stability and resistance of the glass towards crystallization.

The activation energy (Ea) for the crystallization processes can be evaluated from

constant-rate heating DSC scans taken at different heating rates using the Kissinger method

[Kissinger 1957]. The method is based on the assumption that if a reaction proceeds at a

rate varying with temperature i.e. possesses activation energy, the position of the

calorimetric DSC peak, Tp, varies with the heating rate if the other experimental conditions

are maintained fixed [Kissinger 1957]. The variation of the peak temperature can be used

to determine the activation energy of the reaction. According to the Kissinger method, the

crystallization peak temperature, Tp, in the DSC scan depends on the heating rate, φ, as

follows:

⎟⎟⎠

⎞⎜⎜⎝

⎛−=⎟

⎟⎠

⎞⎜⎜⎝

⎛

P

a

P RTE

T 2ln φ+ Constant . (1.10)

By plotting ln(φ/T2P) versus (1/TP), a straight line with slope Ea / R is obtained, where Ea is

the crystallization activation energy and R is the gas constant.

On the other hand, kinetic data on first-order transformations are often obtained by

isothermal analysis. One of the legacies of the classic work done by Kolmogorov

20

[Kolmogorov 1937], Johnson and Mehl [Johnson 1939] and Avrami [Avrami 1939, 1940,

1941] concerning the kinetics of phase transformations involving nucleation and growth

under isothermal conditions is the Johnson-Mehl-Avrami (JMA) [Yavari 1999]

transformation equation.

The crystallized volume fraction during isothermal annealing can be determined

accurately by measuring the partial area of the exothermic signal [Málek 2000, Peng 2005]

assuming that the volume fraction of the transformed material (X) at any given time (t) is

directly proportional to the fractional area of the exothermic peak [Scudino 2008,

Venkataraman 2007b]. According to the JMA equation [Christian 2002b]:

[ ]( )nT tKtX )(exp1)( τ−−−= , (1.11)

where X is the volume fraction of the crystallized phase(s), t the annealing time, KT is a

kinetic constant dependent on the temperature, n is the Avrami exponent and τ the

incubation time for the process. The incubation time is the time period that must elapse

prior to formation of nuclei.

The kinetic constant KT (which can be used to estimate the activation energy for the

transformation, e.g. the devitrification process) is a function of the annealing temperature

and, assuming it to be described by an Arrhenius-type equation, it can be written as

⎟⎠⎞

⎜⎝⎛ −

=RTEKK A

T exp0 , (1.12)

where K0 is a constant and EA is the activation energy for crystallization. The most

important use of this equation has been in the determination of the Avrami exponent. The

Avrami exponent n can vary from 1 to 4 and it is used to describe the transformation

mechanism, such as the nucleation and growth behavior [Christian 2002a]. Based on the n

values, valuable information about the phase transformation can be obtained especially

regarding the nucleation and growth processes as a function of time [Christian 2002b].

Avrami exponent greater than 2.5 implies increasing nucleation rate of all shapes growing

from small dimensions. Constant nucleation rate occurs when n is 2.5 while decreasing

nucleation rate takes place when n is in between 1.5 and 2.5. Zero nucleation rate occurs

when n is 1.5. When Avrami exponent n is 1, it represents formation of needles and plates

of finite long dimensions and thickening of long needles.

21

The values of KT and n can be calculated by using the relation [Avrami 1939, 1940,

1941]

)ln(ln1

1lnln τ−+=⎥⎦

⎤⎢⎣

⎡⎟⎠⎞

⎜⎝⎛

−tnKn

X T . (1.13)

Plotting ln[ln(1/(1-X))] against ln(t-τ) for different annealing temperatures, JMA plots are

obtained. The Avrami exponent n is the slope of the JMA plots.

1.3 Powder metallurgy

Powder metallurgy (P/M) may be defined as a near-net or net-shape manufacturing

process that combines the features of shape-making technology for powder compaction

with the development of the desired microstructures and properties (physical and

mechanical) through subsequent densification or consolidation processes (e.g., sintering)

[Sanderow 1997]. In this process, parts can be produced from metal powders without

passing through the molten state [Dowson 1990]. This process is highly cost effective in

producing simple or complex parts at, or close to, final dimensions with respect to other

fabrication methods like casting, stamping or machining. P/M is the best choice when

requirements for strength, wear resistance or high operating temperatures exceed the

capabilities of die casting alloys [Upadhyaya 2002]. In addition, P/M offers greater

precision, eliminating most or all of the finish machining operations required for castings,

and it avoids casting defects such as blow holes, shrinkage and inclusions [Upadhyaya

2002]. This method is used for several groups of important materials, such as refractory,

composite, porous and glassy materials [Dowson 1990].

P/M processing provides the following advantages over other processing routes

[Trudel 1998]:

Production of complex shapes to very close dimensional tolerances, with minimum

scrap loss and fewer secondary machining operations.

Physical and mechanical properties of the components can be tailored through close

control of starting materials and process parameters.

Particular properties can be improved through secondary processing operations

such as heat treating and cold/hot forming.

22

The main steps in P/M are the powder production and its consolidation. The

sequence of operations to obtain the final product is threefold [Dowson 1990]:

Mixing or blending the powders according to the desired composition and structure.

Loading the powder mix or blend into a suitable die or mould followed by the

powder consolidation through the application of pressure with or without heat in

either controlled or open atmosphere.

Heating the compacts, generally in protective atmosphere, to cause the particles to

bond together. This process is called sintering. The sintering temperature is

normally below, and many cases significantly below, the melting point of the metal

or alloy compact.

1.3.1 Powder production

There are numerous methods of powder production [Dowson 1990], which involve

mechanical (e.g. ball milling) and chemical methods (e.g. reduction of oxides). Mechanical

methods for powder production consist of mechanical comminution, milling and grinding.

Powders produced by chemical methods involve chemical reduction and decomposition of

compounds. Another procedure for powder production is atomization, in which a stream of

molten metal is forced through a small nozzle and then, depending upon the metal involved,

is disintegrated by a jet of water or gas [Dunkley 1986]. The selection of the processing

route for metal powder production is based on the raw material available and the desired

end product and its structure [Upadhyaya 2002]. Suitable methods for powder production

depend on required production rates, powder properties, and the physical and chemical

properties of the material. Chemical and electrolytic methods are useful for producing

high-purity powders [Trudel 1998]. Mechanical milling is the most widely used method of

powder production for hard metals and oxides [Suryanarayana 2001]. Additional milling of

atomized or electrolytic powders is also a very common and economical practice to

produce uniform particle size and shape [Trudel 1998]. Among the different processing

routes, atomization is perhaps the most versatile method that produces metal powders over

a wide range of production rates (from 1 to 105 tons/yr) and a wide variety of powder sizes

from 10 to 1000 µm [Yule 1994].

The most important parameters of any particulate product are particle shape and

size, particle size distribution, purity and apparent density [Dowson 1990]. These

parameters are strictly linked to the powder production process used and they significantly

23

influence the properties of the final material. For example, large particle sizes give rise to

porous end products compared to fine particles. Therefore, a fine grain size is generally

preferred as it also gives greater particle strength, which helps to prevent particle fracture

during compacting [Hirschhorn 1969]. Particle shape is also an important factor as

irregularly shaped particles are required to ensure that the as-pressed component has a high

green strength from the interlocking and plastic deformation of individual particles with

their neighbors. In addition, powder densification can also be remarkably influenced by

particle size distribution [Ferguson 1998].

1.3.2 Metal powder compaction

The compaction of metallic powders has two major functions: to consolidate the

metal powders into desired shape and to sinter the compacts to obtain desired structure and

density [Upadhyaya 2002]. Die compaction process is the one of the most used compaction

method. In this method, shown schematically in Figure 1.6, a die cavity of the desired

shape is filled with the metal powder. Pressure is applied by the axial movement of one or

both punches. The pressure causes the metal particles to mechanically interlock and cold

weld together into a porous mass of the approximate shape and dimensions desired for the

final component. This as-pressed shape, commonly referred to as a green compact, is then

heated to elevated temperatures to achieve full density [Ferguson 1998].

Although most of the sintered parts are made by pressing the powder mix at

ambient temperature followed by sintering, hot pressing is used in certain cases, such as for

hard and brittle materials, where pressure and heat are applied simultaneously. At elevated

temperatures metals are softer and, therefore, it is usually possible to achieve higher

density without increasing the applied pressure [Dowson 1990]. Hot pressing is a suitable

method for densifying materials with poor sintering behavior. This technique, which

combines powder pressing and sintering into one single operation, offers many advantages

over conventional powder consolidation. By the simultaneous application of temperature

and pressure, it is feasible to achieve near theoretical density in a wide range of hard-to-

work materials. As the resistance of metal particles to plastic deformation decreases

rapidly with increasing temperature, much lower pressures are required for consolidation

by hot pressing. In addition, densification by hot pressing is relatively less sensitive to

powder characteristics, shape, size and size distribution, which are critical in cold pressing

and sintering [Upadhyaya 2002].

24

Figure 1.6 Schematic diagram of cross sectional view of uni-axial hot pressing. The die containing the powder is externally heated while pressure is applied through the upper and lower punches (after [German 1996]).

The parameters controlling hot pressing (pressure, temperature, time and the

working atmosphere) largely control the properties of the compacts. The various steps

involved in the hot pressing procedure are the following [Upadhyaya 2002]:

1. Powder or a cold compacted preform is placed into the die mould.

2. The mould is heated either by resistance or by induction to a predetermined temperature.

3. The powder in the die cavity is then pressurized.

4. The temperature is steadily increased during compacting until a maximum required

temperature is reached.

5. Compacting pressure and temperature are maintained for a dwell time and

6. The mould is cooled slowly, under pressure, to a temperature at which oxidation of the

material would not occur.

There are many variations on the general procedure given above. In many cases, it

is preferable to apply a nominal pressure or even the maximum required compacting

pressure before the initiation of consolidation cycle. In place of inert gas, vacuum for hot

pressing offers additional advantages of removing air from the powder body thus

eliminating the possibility of air entrapment [Upadhyaya 2002]. The use of elevated

temperatures and long dwell times allows densities of >95% to be achieved at compaction

pressures that are one third to one half those needed for cold pressing to lower density

levels [Ferguson 1998]. Higher densities can be achieved by extrusion. Extrusion is a

plastic deformation process to produce highly dense bulk samples in which a pre-

25

compacted sample or a billet is forced to flow by compression through the die orifice of a

smaller cross-sectional area than that of the original sample (Figure 1.7).

 

Figure 1.7 Schematic illustration of a vertical uni-axial hot extrusion process.

Because of the large forces required in extrusion, most metals are hot extruded

under conditions where the deformation resistance of the metal is low. Depending on the

material being extruded, hot extrusion is done at relatively high temperatures. Besides the

working temperature, other important parameters of hot extrusion are the extrusion ratio,

the working pressure, the speed of deformation and the frictional conditions and

lubrication. Among these parameters, the extrusion ratio (the ratio of the initial cross-

sectional area of the sample to the final cross-sectional area after extrusion) is the main

factor for achieving the desired density.

1.3.3 Sintering

The ISO (International Organization for Standardization) definition of sintering is:

“The thermal treatment of a powder or compact at a temperature below the melting point

of the main constituent, for the purpose of increasing its strength by bonding together of

the particles” [Dowson 1990]. Bonding together of the particles implies the formation of

bonds in the areas where neighboring particles are deformed at their points of contact by

the applied pressure. During sintering these areas of metallurgical contact grow and the

strength of the sintered body progressively increases [Dowson 1990].

Sintering is a complex process and for any given metal and set of sintering

conditions there are different stages, driving forces and material transport mechanisms

26

associated with the process. The various stages of sintering can be grouped in the

following sequence [Exner 1979, Hirschhorn 1969, Upadhyaya 2002]:

(1) Initial bonding among particles

(2) Neck growth

(3) Pore channel closure

(4) Pore rounding

(5) Densification or pore shrinkage

(6) Pore coarsening

Figure 1.8 Schematic illustration of two-particle model for initial stage of sintering (a) without shrinkage (b) with shrinkage (after [Exner 1979, Kang 2005]).

Bonding takes place very early in the sintering process as the materials is heated up. The

bonding process involves diffusion of atoms leading to the development of grain

boundaries. This takes place at sites where intimate physical contact between adjacent

particles occurs.

Neck growth is the second stage of sintering and is closely related to the first stage of

initial bonding. The newly formed bonded areas are termed necks, which grow in the

second stage of sintering. The neck growth requires the transport of materials within the

sintered mass but does not imply any decrease in the amount of porosity, i.e. no shrinkage

of the material [Exner 1979, German 1996]. Neck growth takes place rather rapidly in the

early stages of the sintering process and continues in the following stages. Figure 1.8

shows the schematic illustration of two-particles sintering together and defining geometry

of neck growth. Initially, the contacts between the spheres are point contacts. After some

27

sintering, due to necking, the contacts become more planar in nature. Neck growth also

results in growth of the initial grain boundaries associated with stage one.

Pore channel closure represents a rather major change in the nature of the porosity in the

sintered mass. Closing off of the tortuous and interconnected pore channels leads to the

development of isolated or closed porosity. Pore channel closure stage may proceed for

some time and overlap stages four and five. The change from interconnected to isolated

porosity can usually be observed microscopically. In particular it is noted that: (a) with

porosities greater than about 10 vol.% most of the porosity is in an interconnected form;

(b) with porosities less than about 5 to 10 vol.% most of the porosity is of the closed or

isolated type [German 1996].

Figure 1.9 Schematic illustration of three particle sintering model: (a) original point contacts, (b) neck growth, (c) and (d) pore rounding (after [Exner 1979, Hirschhorn 1969]).

Pore rounding may be considered as natural consequence of neck growth. When material

is transported to the neck regions from the pore surfaces, the pores themselves become

more rounded as shown in Figure 1.9. With sufficient time at temperature is possible to

achieve almost perfectly spherical pores. Pore rounding is promoted by high sintering

temperatures. This stage of sintering is particularly important with respect to the influence

of porosity on the mechanical properties of the sintered materials [Exner 1979, German

1996].

Pore shrinkage and eventual pore elimination is often considered as the most important

stage of sintering. Only with sufficient time at temperature may it evidence itself by

densification of the sintered mass. It is important to realize that the process of pore

28

shrinkage, leading to a decrease in the volume of the sinter mass, must involve movement

of the solid into the porosity.

Pore coarsening usually takes place after most of the other stages of sintering have

occurred. The process simply consists of the shrinkage and elimination of small isolated

pores and the growth of larger ones. The total amount of porosity associated with all these

pores remains the same, but the number of pores decreases and the average size increases.

Hence, no densification of the material is associated with this stage [German 2008].

Driving force for sintering

At elevated temperatures, the loose powder or the compact is not at equilibrium and

it is prone to substantial changes in its internal structure towards a more stable state [Tilley

2004]. The driving force for the change is the reduction in free energy of the system. In the

sintering process, the necessary reduction of free energy is associated with the decrease in

internal surface area of the sintered mass. A decrease in surface area corresponds to a

decrease in the surface free energy contribution to the total free energy of the system (i.e.

area multiplied by specific surface or interfacial free energy). The total free energy of a

powder compact is expressed as γA, where γ is the specific surface (interface) energy and A

the total surface (interface) area of the compact. The variation of the total free energy can

be expressed as [Exner 1979, Kang 2005]:

AAA Δ+Δ=Δ γγγ )( . (1.14)

Here, the change in interfacial energy (Δγ) is due to densification and the change in

interfacial area ΔA is due to grain coarsening. For solid state sintering, Δγ is related to the

replacement of solid/vapor interfaces (surface) by solid/solid interfaces [Kang 2005].

In other words, the sintered mass undergoes changes that tend to eliminate the

internal surface area. For example, pore rounding reduces the surface area while

maintaining the amount of porosity at a constant level. This is because the ratio of surface

area to volume is reduced when the shape of the pore approaches a sphere. The surface

area can then be reduced by the pore shrinkage and the densification stage of sintering. The

ratio of surface area to volume is also decreased by increasing the average size of the pores

while maintaining the total volume constant (i.e. pore coarsening). As well, pore channel

closure decreases the surface/volume ratio. Hence, the greater the amount of surface area

in the original materials, the greater the driving force for sintering [Tilley 2004].

29

Transport mechanisms

Sintering can be considered as “a thermally activated material transport in a powder

mass or a porous compact, decreasing the specific surface by growth of particle contacts,

shrinkage of pore volume and change of pore geometry” [Thümmler 1993]. Accordingly,

most sintering theories are based on transport phenomena associated with a particular stage

of sintering. In the following, the major mechanisms of material transport are presented

[Kang 2005]. Different sintering mechanisms are illustrated in two-particle model in the

Figure 1.10

Figure 1.10 Schematic illustration of different material transport paths during sintering as applied to the two-sphere sintering model (E-C, evaporation-condensation; SD, surface diffusion; VD, volume diffusion; GBD, grain boundary diffusion;) (after [Exner 1979, Kang 2005, Schatt 1987, Schatt 1985a, Schatt 1985b]).

Evaporation and condensation: because of the higher vapor pressure over convex surfaces

as compared to the neck regions, it is possible in some systems for material to be

transported as vapor to the neck region. Neck growth, pore rounding and pore channel

closure can be accomplished by this mechanism. This mechanism is important for

materials with relatively high vapor pressure, so that significant amount of material can be

transported [Exner 1979, German 1996].

Volume (lattice) diffusion: This transport mechanism is widely accepted for the sintering of

metallic materials. Volume or lattice diffusion refers to the movement of atoms within the

solid crystalline material. The most prevalent specific type of atomic motion is the

“vacancy exchange” mechanism. This process involves the movement of atoms into vacant

lattice sites (i.e. vacancies). If there is a directionality associated with a substantial amount

30

of such atomic motion, then there is a net transport of material in a specific direction.

Diffusion of atoms along a specific direction (rather than a random motion) is a

consequence of chemical potential gradients existing within the solid. Material transport by

volume diffusion is due to the existence of vacancy concentration potentials (differences)

in the solid, and the movement of vacancies from regions of high concentration to regions

of low vacancy concentration [Brand 1993, German 1996].

Surface diffusion: Atomic transport in the solid state can also occur by surface diffusion, i.e.

the motion of atoms on external surface. In analogy with volume diffusion, the most

probable mechanism of surface diffusion is the exchange between surface atoms and

surface vacancies. This process is particularly important in the first sintering stages, when

the specific surface is still high. It is generally accepted that surface diffusion does not

cause pore shrinkage and hence densification, however, it can promote neck growth

[Hirschhorn 1969].

Creep deformation: This process, which occurs at elevated temperatures and under

constant load, is a consequence of the repetition of the following steps: (1) generation of

dislocations activated by the applied stress, their movement and their arrest at some

obstacles; (2) dislocation climb and generation of new dislocations. The rate controlling

step is the diffusion dependent climb. There is a stress concentration at the head of the

dislocation pile-up which can lead to the production or annihilation of vacancies. The

driving force for this material transport mechanism is the presence of shear stresses in the

solid. Those factors that promote stresses in the solid, such as a large curvature in the neck,

small pore radius and high surface tension, would tend to increase the probability of having

this mechanism be the dominant one for sintering [German 1996, Hirschhorn 1969, Schatt

1983].

Transformations during sintering

In this section, the main changes or transformations that may take place during

sintering are considered [Hirschhorn 1969].

Grain growth: Grain growth is the most important transformation that occurs during

sintering. In normal sinter mass there is a very large amount of grain boundary areas

because of the small sized particles. Grain boundaries represent a positive contribution to

the free energy of the material; therefore, a large driving force exists for removing the

grain boundaries and for grain growth to reach the lowest energy state (higher degree of

stability). This is why there is substantial grain growth during sintering. Almost any

31

deviation from pure single−phase material will reduce the tendency for grain growth. This

includes (i) grain boundary pinning to decrease grain boundary mobility through residual

pores [Hahn 1990], impurities and solutes [Averback 1993] and second phase particles

[Hillert 1988], and (ii) reduction of the driving force for grain growth by lowering the grain

boundary energy through the addition of solute atoms that segregate at the grain

boundaries [Koch 2009].

Alloying: Very often the sintering operation is used to produce homogeneous alloys from

the original mixture of two or more elemental powders. Alloying during sintering is due to

diffusion. Although most of the available experimental evidence indicates that volume

(lattice) diffusion is the main mechanism for alloying, surface diffusion might become the

dominant mechanism, particularly at low temperatures.

Phase transformations: Many types of phase transformations may occur in the solid state

during sintering at a constant temperature or during the cooling of the material from the

sintering temperature. In some cases such transformations would follow sufficient alloying.

Probably the best example of a phase transformation associated with sintering is the

production of sintered steels. Plain carbon steels would be made by mixing graphite and

iron powder; during sintering the iron and graphite would alloy to form the high

temperature austenite phase (a solid solution of carbon in fcc iron). Once the austenite is

formed then the desired pearlitic structure can be obtained upon cooling.

Influence of material and process parameters

The major variables which determine sinterability and the sintered microstructure

of a powder compact may be divided into two categories: material variables and process

variables. The variables related to raw materials (material variables) include chemical

composition of powder compact, powder size, powder shape, powder size distribution,

degree of powder agglomeration, etc. These variables influence the powder compressibility

and sinterability (densification and grain growth). The process variables involved in

sintering are mostly thermodynamic variables, such as temperature, time, atmosphere,

pressure, heating and cooling rate [Exner 1979, German 1996, Kang 2005, Schatt 2007].

Particle size: Decreasing particle size leads to improved sintering. With a smaller particle

size there would be greater inter-particle contact (number of necks) and, hence, more paths

for volume diffusion. Also, a small particles size may correspond to a smaller grain size,

promoting transport by grain boundary diffusion.

32

Particle shape and surface morphology: The optimization of these parameters may lead to

improved intimate physical contact among particles in the sinter mass and increased

internal surface area, promoting sintering. For example, irregularly shaped particles lead to

higher density of the compacts than spherical powders. As well, increasing micro and

macro-surface roughness may assist sintering.

Particle structure: A fine grain structure of the original particles can promote sintering

because of its favorable effect on several material transport mechanisms. The presence of a

large amount of lattice imperfections, such as dislocations, usually resulting from plastic

deformation, can promote sintering because such defects enhance sintering by diffusional

processes.

Particle composition: The driving force for sintering may be either increased or decreased

by alloying elements or impurities in the material. Diffusional mass transport may also be

affected by the presence of alloying or impurity atoms in the lattice.

Temperature: Increasing the sintering temperature greatly increases the rate and magnitude

of any changes occurring during sintering. Many investigations of sintering have indicated

that the precise material transport mechanism, which controls the rate of sintering,

remarkably changes with varying temperature.

Time: Although the degree of sintering increases with increasing time, the effect is small in

comparison to the temperature dependence. The rate of sintering decreases with increasing

time.

Pressure: Combination of temperature and pressure induce accelerated densification

process and elimination of residual pores. Pressure assisted sintering adds stress to

accelerate the material flow during sintering [German 1996].

33

Chapter 2: Sample preparation and characterization

2.1 Preparation of amorphous and partially amorphous powders

The samples investigated in this work were produced through two ways: solid-to-

solid transformation, i.e. by mechanical alloying of elemental powder mixtures, and

through liquid-to-solid transformation, i.e. by melt spinning and gas atomization.

The starting material for mechanical alloying consisted of a mixture of highly pure

elements (Table 2.1) of the given nominal compositions. Differently, pre-alloyed materials

were used as starting materials for the melt spinning experiments. The pre-alloyed

materials were prepared by mixing the appropriate weights of each chemical constituent, in

the form of small lumps, for the desired atomic composition. The lumps were mechanically

cleaned to remove any possible surface oxide layer. The elements were then melted in an

induction furnace in an argon gas atmosphere purified with titanium getter. The rods were

remelted several times to ensure homogeneity in composition. The molten alloy was then

cast into cylindrical rods with 10 mm diameter and 100 mm length by copper mold casting

under argon atmosphere.

Table 2.1. Starting materials for mechanical alloying experiments.

Element Supplier Purity Particle size ≤

Al Alfa Aesar 99.5% 45 µm

Y Mateck 99.9% 420 µm

Ni Alfa Aesar 99.5% 300 µm

Co Alfa Aesar 99.9% 300 µm

2.1.1 Melt spinning

Glassy ribbons were prepared in a single-roller melt spinner (Edmund Bühler D-

7400) under argon atmosphere. In this apparatus (Figure 2.1) a piece of the pre-alloyed rod

(5-7 g) was molten inductively in a quartz tube having a rectangular slit, at the end of the

nozzle, of 3 mm length and 0.7 mm width. The position of the nozzle tip can be adjusted

with respect to the wheel surface so that the molten alloy was perpendicularly ejected onto

the wheel surface from a distance of about 0.3 mm. The whole unit is enclosed in a

34

chamber connected with vacuum pumps and an argon inlet. The chamber was evacuated to

10-3 Pa and rinsed with argon two times. The temperature of the liquid metal was

monitored by an optical pyrometer and when the temperature rises about 150-200 K above

the melting point of the alloy (typically 1200-1400 K), an overpressure of about 250 mbar

of pure argon was applied from

   

Figure 2.1 (a) Schematic diagram of melt spinning process [Suryanarayana 2001]

(b) single roller Edmund Bühler melt spinning device.

an external reservoir to eject the molten alloy out of the quartz tube onto the external

surface of a copper wheel rotating at a velocity of 40 ms-1. The melt-spun ribbon, which

detached from the wheel surface during melt spinning, was guided towards a collecting

box. The resulting ribbons had typical widths of 3 mm and thicknesses of 40-50 µm.

2.1.2. Ball milling

Milling experiments, starting from elemental powder mixtures with nominal

composition Al85Y8Ni5Co2, were performed using a Retsch PM400 planetary ball mill and

hardened steel balls and vials (Figure 2.2(a)). The schematic illustration of ball milling is

shown in Figure 2.2(b). In this type of mill, four vials are arranged eccentrically on the

supporting disc (sun wheel) of the planetary mill. While operating, the rotation of the

supporting disc is accompanied by a rotation of the vials around their own axes, in the

opposite direction. As a result of the superposition of opposite rotating motions of

supporting disc and vials, the steel balls in the grinding vials are subjected to superimposed

rotational movements, the so-called Coriolis forces. The difference in speeds between the

balls and grinding vials produces an interaction between frictional and impact forces,

35

which releases high energies. The rotational velocity of the supporting disc can be

considered as a rough estimate of the milling intensity. This can be controlled and kept

constant at values ranging from 60 to 400 revolutions per minute (rpm). In the present

work, the rotational velocity was set to 150 rpm for the milling experiments. Milling was

carried out as a sequence of 15 min milling intervals interrupted by 15 min breaks to avoid

a strong temperature rise until no reflections from the starting metallic elements were

detected in the X-ray diffraction pattern. The starting materials for the milling experiments

(typically 30 g) were charged in the milling vials equipped with a flexible “O”-ring,

together with 10 mm-diameter steel balls to give a ball-to-powder mass ratio (BPR) of 10:1.

To avoid any possible atmosphere contamination during milling, vial charging and any

subsequent sample handling was carried out in a Braun MB 150B-G glove box under

purified argon atmosphere (less than 1 ppm of O2 and H2O).

Figure 2.2 (a) Retsch PM400 planetary ball mill chamber showing the vials and balls and (b) schematic illustration of the ball milling principle [Suryanarayana 2001].

Additionally, Al-based glassy powders were produced by controlled milling at

cryogenic temperature of melt-spun glassy ribbons with nominal composition

Al85Y8Ni5Co2. The ribbons were milled for 5 h at a ball-to-powder mass ratio of 10:1 and

at rotational velocity of 150 rpm. The milling was carried out with a BPR of 10:1 as a

sequence of 15 minutes milling intervals interrupted by 15 minutes breaks.

2.1.3 Gas atomization

Gas atomization is one of the most used processing routes for the production of

metallic powders due to its advantages, such as homogeneous and uniform size distribution

36

of particles and possibility to scale up the process to tonnage quantities. Figure 2.3 shows a

schematic representation of a gas atomizer. Atomization consists of breaking up of bulk

liquid metal or alloy into fine droplets and allowing them to solidify as powder particles

[Lawley 1977]. Mostly, these particles are round due to the liquid surface tension. This

property causes thin ligaments of liquid to be unstable; that is, they break up into droplets,

or atomize. The main fluid properties which affect the size of the droplets are surface

tension, viscosity and density. The average droplet size is larger with higher density,

viscosity and surface tension [Lawley 1978]. Atomization is a popular route for large scale

powder production. There are different atomization processes in industrial practice to

produce different metal or alloy powders. Most common atomization processes can be

mainly classified into gas, water and centrifugal atomization processes [Lawley 1978]. In

gas atomization a continuous stream of liquid metal is broken down into droplets by means

of a subsonic or supersonic gas jet. Atomization occurs by the kinetic energy of the

atomizing medium, typically nitrogen, argon, or air. Various atomization geometries are

used in commercial practice. The process is governed by a number of interrelated operating

parameters. Controllable variables include jet distance, jet pressure, nozzle geometry,

velocity of gas and metal, metal superheat, angle of impingement, metal surface tension

and the melting range in case of alloy powder production [Lawley 1978].

Figure 2.3 Schematic representation of a gas atomization unit ([German 1984]).

37

In gas atomization, the nozzle geometry is a critical aspect to govern particle size,

shape, distribution and yield of the powder [Dunkley 1986]. The solidification rate in gas

atomization depends on the partic

Co

GAP were prepared by high pressure Ar ga

5

m

pure elem

le size, with higher rates (104-106 K/s) associated with

smaller particle sizes (~20-40 μm) [Lawley 1977], as well as on the type of the atomization

medium. Higher solidification rates are achieved with smaller particle sizes and lighter

gases. Additionally, it is possible to increase the solidification rate by cooling the

atomizing gas, or even by pumping additional pre-cooled gas from the bottom of the

chamber. If sufficient superheating is provided to the alloy and the atmosphere being

neutral, the final powder product is spherical. The range of powder sizes is broad with the

mean particle diameter of around 100 μm, even though a mean diameter of 12 to 15 μm

has been also reported [Lawley 1977]. Typically, the yield of powders in gas atomization

can go up to as high as 80% of the starting material. Typical solidification rates achievable

are 102- 106 K/s, especially for powder particles smaller than 30 μm [Lawley 1977].

In this work, Al-based gas-atomized powders (GAP) with different compositions

have been obtained from different sources. The Al Gd Ni Co and Al Y Ni84 6 7 3 90.4 4.3 4.4 0.9

s atomization at the Materials Processing

Center, Ames Laboratory, Ames (USA). This was achieved by mixing the appropriate

weights of each chemical constituent in the form of small lumps, for obtaining the desired

composition and heating them in a graphite crucible. The elements used had purities

ranging from 99.9% to 99.999%. Prior to heating, the lumps were mechanically cleaned to

remove any possible surface oxide layer. Subsequently, the melt was cast in a water-chilled

Cu mold. The powders were then produced by high pressure Ar gas atomization using a

close coupled annular nozzle having a melt delivery inner diameter of 3.2 mm. Following

atomization, the powders were screened using sieves to a size below 106 μm. The particle

size analysis reveals that the median size (X50) of the as-atomized powders Al84Gd6Ni7Co3

and Al90.4Y4.3Ni4.4Co0.9 are 23 µm and 21 µm respectively.

The Al87Ni8La GAP were prepared using the “Nanoval” process [Gerking 1993] at

NANOVAL GmbH & Co. KG, Berlin (Germany). The master alloy was prepared fro

ents Al (99.999%), Ni (99.95%) and La (99.5%) by induction melting in a high

frequency electromagnetic levitation furnace under a purified argon atmosphere. The

levitated melt was kept at high temperatures until the surface oxide skin broke up. After 2

to 7 min of holding, the high frequency current was switched off and the levitated melt fell

down into a water chilled copper crucible, where it solidified forming a crystalline ingot.

This procedure was repeated several times to ensure the homogeneity of the alloy. The

38

melt of the Al87Ni8La5 alloy was then heated up to approximately 1350 K, and was

subsequently gas-atomized using argon gas. The gas-pressure was controlled manually and

could reach a maximum value of 30 bar. Atomization yielded two powder fractions: a

coarser major powder fraction collected in the atomizer recipient and a finer minor powder

fraction collected in the cyclone. A slow oxidation of the as-atomized powder was allowed

by flooding the atomizer with air.

2.2 Powder consolidation

methods, such as hot pressing, hot extrusion and spark

plasma sintering were used in this work to consolidate the powders produced by milling

.2.1 Hot pressing and extrusion

Uni-axial hot pressing is one of the simplest powder consolidation processes where

to the powder placed in a die–punch setup at high

temper

ity. An electro-hydraulic universal axial pressing

machin

Different consolidation

and gas atomization processes into highly dense bulk samples.

2

high pressure is applied uniaxially

ature to induce particle sintering.

The use of elevated temperatures and pressures along with long dwell times allows

densities of > 95% of the theoretical dens

e made by WEBER PWV 30 EDS, Germany, with a capacity of 350 kN maximum

load was used to consolidate the powders by hot pressing. Before starting hot pressing all

the parts such as compaction die and punches are cleaned and sprayed with a thin layer of

boron nitride for lubrication purpose. The whole setup is placed inside the closed chamber.

Approximately 2 to 3 g of powder is placed in the 10 mm diameter of die. The temperature

was measured by a thermocouple of Pr/Rh Pt which was fixed in a dedicated cavity within

the die, ensuring a continuous monitoring of the operating temperature. The chamber is

evacuated to 1×10-3 Pa before starting hot pressing for degassing and to minimize the

oxidation during hot pressing. Careful selection of pressure, temperature and dwell time

has been chosen according to the desired microstructure after hot pressing. Desired

pressure (typically 500 MPa) is applied, then die and punch setup is heated to the desired

temperature with an inductive coil. The compaction temperature and dwell time was

chosen from the thermal studies performed using differential scanning calorimeter. After

39

finishing the heating cycle the chamber is left in vacuum to cool down and Argon was

purged to remove the sample from the chamber. Full density is usually not achieved, and 2

to 4% porosity remains in consolidated sample. For this reason hot-pressed samples are

often hot extruded to obtain highly dense samples.

Hot extrusion is a process by which the billet or pre-compacted samples are forced

through the die cavity to obtain long straight samples by applying high pressure (about

500 M

2.2.2 Spark plasma sintering

Spark Plasma Sintering (SPS) is also known as pressure-assisted resistance or

charge powder compaction, electroconsolidation, plasma

tivate

Pa) as well as high temperature to obtain highly dense samples. The hot extrusion is

performed using the same equipment which is used for hot pressing with a typical

reduction of the sample diameter from 10 to 4 mm (extrusion ratio ~ 6:1). In order to

facilitate easy flow of the material during extrusion, copper jackets were used for the

extrusion of the samples pre-compacted by hot pressing.

electric-discharge sintering, dis

ac d sintering, field-assisted sintering, electric pulse sintering, and pulse electric

current sintering. The schematic diagram of a typical SPS machine is shown in Figure

2.4(a).

   

Figure 2.4 (a) Schematic illustration of SPS and (b) Picture of Dr. Sinter 515S SPS

machine.

ark plasma sintering is a newly developed processing technique which makes

possible sintering and sinter-bonding at low temperatures and in short periods by high

energy

Sp

pulsed electric energy [Tokita 1993]. SPS of metallic powders involves the passage

of pulsed electric current through the powder particles while subjected to an applied

40

pressure. The current pulses generate Joule heating and an electric field diffusion effect. In

the SPS process, the powder particle surfaces are more easily purified and activated than in

conventional electrical sintering processes and material transfers at both the micro and

macro levels are promoted, so a high-quality sintered compact is obtained at a lower

temperature and in a shorter time than with conventional sintering processes [Tokita 1999].

Advantages of the process over conventional hot pressing include sintering at significantly

lower temperatures, shorter sintering times, higher heating rates, accurate control of

sintering energy, high reproducibility and much faster phase transformation kinetics than

conventionally possible [Garay 2003]. The SPS process is widely used to produce

functionally graded materials, intermetallic compounds, nanocrystalline materials, metallic

glass materials, metal matrix composites and hard-to-sinter materials.

In this work, consolidation by SPS was carried out at the Research Center for

Machine Parts and Materials Processing, University of Ulsan, Republic of Korea using a

DR.SIN

2.3. Sample characterization

lidated samples prepared by the different techniques

were investigated using several analytical methods. The phases and the microstructure

wer c

TER SPS machine SPS-515S with 50 kN maximum load (shown in Figure 2.4 (b)).

The system consists of a SPS sintering machine with a vertical single-axis pressurization

mechanism, specially designed punch electrodes incorporating water cooler, a water-

cooled vacuum chamber, a vacuum/air/argon-gas atmosphere control mechanism, a special

DC-pulse sintering power generator, a cooling water control unit, a position measuring unit,

a temperature measuring unit, an applied pressure display unit and various interlock safety

units [Tokita 1997]. The powder to be consolidated was placed in WC dies with 10 mm

diameter and then a pressure of about 500 MPa was applied. The chamber is then

evacuated to 10 -3 Pa. SPS generator was switched on to obtain the desired temperature.

The temperature is measured by the thermocouple attached near to the die and punch

assembly. Sufficient dwell time (about 3 min in the current experiments) is maintained to

obtain uniform temperature through out the sample after reaching the desired temperature.

The powders and the conso

e haracterized by X-ray diffraction (XRD), Optical microscopy (OM), scanning

electron microscopy (SEM) and transmission electron microscopy (TEM). The thermal

stability and the crystallization behavior were investigated by differential scanning

41

calorimetry (DSC), whereas the effect of temperature on the viscosity of the powders was

analyzed by parallel plate rheometry.

2.3.1 X-ray diffraction

Standard X-ray diffraction (XRD) in reflection configuration was carried out using

-Brentano diffractometer using Co-Kα radiation (λ = 0.17889 nm),

equ pe

metry

The thermal stability and the crystallization behavior of the samples were

ate heating) and isothermal mode by differential

sca n

nt of intersection

betw

a Philips PW 1050 Bragg

ip d with a secondary graphite monochromator and a sample spinner. The

diffractometer operated at a voltage of 40 kV and a current of 40 mA. The diffracted

intensities were recorded between 20 and 100 degrees (2θ) in a step mode with a step size

of Δ(2θ) = 0.05° and typical counting times between 15 and 30 s per step, longer times

being required for amorphous materials and for samples characterized by particularly small

grain size of the precipitates. Additionally, selected samples were investigated by XRD in

transmission configuration using a STOE Stadi P diffractometer (Cu Kα radiation; λ =

0.15406 nm). The Rietveld method was applied for the profile-fitting structure refinement

using the WinPlotR software package [Roisnel 2001]. For powder samples, a resin slightly

diluted with acetone was used to fix the powders on PVC sample holders. Both the fixing

agent and the sample holder did not interfere with the measurements performed in the

scanning range used in this work.

2.3.2 Differential scanning calori

investigated in both isochronal (constant-r

nni g calorimetry (DSC) using a computer-controlled Perkin-Elmer DSC7 under a

continuous flow of purified argon. Alumina crucibles were used as sample holders. They

were charged with 10 to 20 mg of material. The calibration of the DSC system for all the

heating rates utilized was done using zinc and indium standard samples.

The isochronal DSC studies were done at different heating rates in order to determine

the onset temperature of the glass transition, Tg estimated as the poi

een the linearly extrapolated curve below the transition with the steepest tangent of

the rise in the heat flow signal, the crystallization temperatures, Tx determined as the onset

temperature of the exothermic events with peak temperature, TP and the crystallization

enthalpies, ΔHx corresponding to the areas of the exothermic crystallization events. The

glass transition is observed as a characteristic endothermic event, whereas crystallization is

42

associated with one or more sharp exothermic peaks. For each individual sample, two

successive DSC isochronal runs were recorded at the selected heating rate followed by

cooling down to room temperature at 100 K/min. The second run of the specimen served as

a baseline. Subtraction of this baseline from the first run realized the correction for the

apparatus specific baseline shift.

The isothermal DSC studies were carried out at the annealing temperatures Ta < Tx

and were carried out in order to investigate the kinetics of the phase transformations.

Sim

The viscosity of the glassy powders as a function of temperature was measured

eometry using a computer controlled thermal mechanical

ana ze

ilarly to the isochronal scans, two runs were recorded for each sample. The second run

was used as baseline for the data from the first run.

2.3.3 Parallel plate rheometry

isochronally by parallel plate rh

ly r (Perkin-Elmer TMA7). Measurements on powder samples were done under a

purified argon atmosphere at a heating rate of 20 K/min using quartz penetration probes

(3.7 mm diameter) as parallel plates and a static force of 2.6 N. Before each measurement,

the samples were first heated above the glass transition to the temperature Tx – 20 K in

order to achieve the same relaxed isoconfigurational state and to completely fill the area

between the plates.

Figure 2.5 Derivation of the viscosity from parallel plate rheometry measurements through Stefan’s equation [Busch 1998].

late rheometry measurements. Figure 2.5(a) displays a

cha ct

As a typical example, Figure 2.5 shows how the viscosity of the supercooled liquid

can be evaluated by parallel p

ra eristic isochronal TMA curve, which represents the variation of the height of the

sample as a function of temperature or time. The viscosity η can be derived from the

43

change of the height of the sample versus time through Stefan’s equation [Busch 1998],

where F is the applied load, a is the radius of the plates and h is the height of the sample.

This allows viscosity measurements in the range from 105 to 1010 Pa*s [Busch 1998,

Deledda 2004]. As the glass transition temperature (Tg) is reached and the glassy solid

transforms into the SCL, the curve in Figure 2.5(b) displays a strong viscosity drop. At Tx,

when the crystallization sets in, the viscosity abruptly increases with increasing

temperature, indicating the loss of liquid-like behavior.

2.3.4 Optical microscopy (OM)

Optical microscope, Nikon Epiphot 300 is used to obtain magnified images of the

ble of magnifying the samples from 5 to 100 times with

differe

In this work, SEM characterization of powder and consolidated samples was

ini 1530 (Zeiss) SEM with FEG-Source

(Sc tt

nanometers size. In general, images can be obtained

usin s

n microscopy

TEM studies were performed using a Philips CM20 FEG TEM (USA) operating at

nit (STEM) with bright and dark filed detector,

energy

samples. This microscope is capa

nt objective lenses. This microscope is equipped with in built camera and also

connected with a computer program (a4i Docu from Olympus Deutschland GmbH,

Germany) so that images can be captured by using the software.

2.3.5 Scanning electron microscopy (SEM)

carried out by using a high-resolution Gem

ho ky type), EDS for elemental analysis with Si(Li) detector and QUANTAX

evaluation software (Bruker AXS).

Scanning electron microscopy (SEM) gives high resolution images with details

from millimeters to microns or even

g econdary electrons as well as by back scattered electrons. Back scattered electrons

are normally used to detect contrast between areas with different chemical compositions

while the secondary electron image gives topographical information. Additionally, SEM

can be equipped with an energy dispersive X-ray spectrometer (EDS) or a wavelength

dispersive X-ray spectrometer (WDS). The EDS system provides rapid chemical

composition and mapping.

2.3.6. Transmission electro

200 kV with field emission gun, scanning u

dispersive X-ray spectrometer (EDXS) from Noran Voyager IIa and TV-rate

camera from TVIPS FastScan F114NX. The TEM samples were prepared by mechanical

44

grinding and ion polishing. The samples were ground to a thickness of about 200 μm using

800 and 1200 grit SiC papers. This was followed by dimpling (GATAN dimpler) to

achieve a thickness of about 30 μm at the center of the sample. Final thinning was

performed in a GATAN precision ion polishing system (PIPS), operating at 4-5 kV with an

angle of 2-4°, in an argon atmosphere.

In transmission electron microscopy (TEM), a beam of electrons is transmitted

through an ultra thin specimen, interacting with the specimen as it passes through. An

image

2.4 Mechanical properties

The density measurements of the consolidated samples were performed with a

weighing instrument. The sample was first weighed in air and is weighed in

the eth

The microhardness was measured using a computer-controlled HMV Shimadzu

g machine. The device is equipped with a typical diamond indenter

in the f

is formed from the interaction of the electrons transmitted through the specimen; the

image is magnified and focused onto an imaging device, such as a fluorescent screen, on a

layer of photographic film, or to be detected by a sensor such as a CCD camera. In TEM,

atomic-scale lateral resolution can be achieved without depth information as this technique

works by having the probe electron beam transmitted through a sample that is up to 200

nm thick.

2.4.1 Density

Mettler Toledo

er liquid medium. The humidity and air temperature is also considered in measuring

the density. The density of the sample was then calculated by the Archimedes principle.

2.4.2 Microhardness

Vickers hardness testin

orm of a pyramid with square base and an angle of 136° between the opposite faces.

The applied load was 1.96 N for 10 sec. The diagonal of the imprints as well as the

hardness were calculated using a digital video measuring system. For the indentations, the

samples were embedded in epoxy resin and the measured surface was carefully polished

with a paste containing diamond particles with a diameter smaller than 0.25 µm. The final

result is an average of more than 20 measured data.

45

2.4.3 Compression and tensile testing

The room temperature stress versus strain under compression was measured for

cylindrical samples and bars. The samples for compression testing were cut into cylinders

or cubes with the length equals to twice the diameter or side of a cube, from the hot-

pressed, SPS and extruded samples. Both ends of the specimens were polished to make

them parallel to each other prior to the compression test. WC steel loading platens

lubricated with MoS2 grease were used for the compression tests. The dog-boned shape

tensile specimen with the dimension of gauge length of 20 mm, gauge diameter of 3 mm

and with 4 mm diameter grips were prepared by using a small turning machine. A laser

extensometer-type parallel scanner from Fiedler Optoelektronik GmbH, Germany, was

used for non-contact measurement of strain in tension or compression of specimens under

uniaxial load. The samples were tested in compression or tension with an INSTRON 8562

testing facility under quasistatic loading (strain rate 1×10-4 s-1) at room temperature. A

Bluehill mechanical testing software was used to obtain engineering and true stress – strain

diagrams.

46

Chapter 3: Characterization of Al-based amorphous powders

produced by mechanical alloying and melt spinning

Due to their relatively low glass-forming ability, Al-based amorphous alloys can

generally be obtained only by melt-spinning in the shape of ribbons or by mechanical

alloying (MA) in the form of powder. Accordingly, the first part of this chapter deals with

the production, structural and thermal characterization of mechanically alloyed Al-based

metallic glass with composition Al85Y8Ni5Co2. The second part is devoted to the

production and characterization of Al85Y8Ni5Co2 glassy ribbons produced by melt spinning,

followed by their pulverization under controlled milling conditions in order to produce

particulate Al-based glassy reinforcements for Al-based metal matrix composites (MMCs).

Finally, structural, thermal and mechanical characterization of the glass-reinforced MMCs

are presented together with the modeling of their mechanical properties.

3.1 Mechanically alloyed Al85Y8Ni5Co2 glassy powders

This section focuses on the production of Al85Y8Ni5Co2 glassy powder by

mechanical alloying. The Al85Y8Ni5Co2 alloy has been chosen because of its relatively

good glass-forming ability among the different Al-based glass-forming systems. The

presence of the rare earth element Y and transition element Ni in Al85Y8Ni5Co2 increases

the glass-forming ability of the alloy [Zhang 2007]. In addition, nickel has, among all the

transition metals, the largest negative heat of mixing with aluminum (-22 kJ/mol) [Zhang

2007], which may assist the glass-forming ability of the alloy. Furthermore, addition of

cobalt improves the mechanical properties of the alloy, namely hardness and strength

[Zhang 2007].

Börner et al. [Börner 2001] successfully obtained complete amorphization of the

Al85Y8Ni5Co2 alloy by mechanical alloying and no indications of remaining crystalline

phases were found by X-ray diffraction and TEM. Complete amorphization can be

achieved in this system by employing a systematic variation and optimization of the

milling parameters, i.e. using proper milling conditions such as interval milling at a low

intensity corresponding to a rather low kinetic energy [Börner 2001, Suryanarayana 2001].

Similarly, in this work mechanical alloying of the Al85Y8Ni5Co2 alloy was carried out at

47

low milling speed as a sequence of milling intervals of 15 minutes interrupted by breaks of

15 minutes to avoid a strong temperature rise.

Figure 3.1 XRD patterns (Co-Kα radiation) for the mechanically alloyed Al85Y8Ni5Co2 powder: as-milled, heated to 653 K and heated to 748 K.

Figure 3.1 shows the XRD pattern of the Al85Y8Ni5Co2 powder mechanically

milled for 500 h. The pattern shows the presence of a broad diffuse halo in the range 40 −

50° (2θ), characteristic for amorphous material. Superimposed are broad and low intensity

diffraction peaks, which indicate the presence of a small amount of a crystalline phase,

most likely the Al3Y (space group R-3m).

Figure 3.2 DSC scan (heating rate 20 K/min) for the mechanically alloyed Al85Y8Ni5Co2 powder.

48

Figure 3.2 presents the constant-rate heating DSC scan (20 K/min) of the

mechanically milled powder heated to 823 K. The DSC curve is characterized by the

presence of two sharp exothermic events with onset temperatures Tx1 = 613 K and Tx2 =

669 K, followed at higher temperature by a broad exothermic peak at about 740 K. No

distinct glass transition can be observed.

In order to study the structural evolution during heating, samples of the

Al85Y8Ni5Co2 powder were annealed in the DSC by continuous heating at 20 K/min up to

different temperatures and then were cooled to room temperature at 100 K/min. The phases

formed were identified by X-ray diffraction and their patterns are shown in Figure 3.1. The

XRD pattern after heating up to the completion of the first exothermic DSC peak (653 K)

reveals the diffraction peaks from fcc Al (space group Fm-3m). An overlapping broad

diffraction peak at about 43° (2θ) can also be observed, which reveals the presence of a

residual glassy phase after the first exothermic event. When the sample is heated to 748 K,

i.e. above the second exothermic event, the XRD pattern displays the presence of

diffraction peaks from fcc Al along with the peaks from the intermetallic phases Al3Y,

Al4Ni3 (space group Ia-3d) and Al9Co2 (space group P21/a).

Although the DSC scan in Figure 3.2 shows no clear glass transition prior to

crystallization, glass transition may, nonetheless, occur. A better insight into the flow

behavior of the supercooled liquid can be derived from viscosity measurements [Busch

1998, Deledda 2004]. Accordingly, the influence of temperature on the viscosity of the as-

milled powder was studied by parallel plate rheometry (Figure 3.3).

Figure 3.3 Temperature dependence of the viscosity for the mechanically alloyed Al85Y8Ni5Co2 powder (heating rate 20 K/min).

49

The viscosity η was measured from the change of the height of the sample versus

time according to the Stefan’s equation shown in Figure 2.5. As the glass transition

temperature is reached (Tg1 at about 600 K) and the glassy solid transforms into the super-

cooled liquid (SCL), the curve in Figure 3.3 displays a strong viscosity drop. At 625 K

(Tx1) the alloy starts to crystallize, leading to a strong viscosity increase, which indicates

the loss of liquid-like behavior. A second drop of viscosity (Tg2), probably related to the

glass transition of the residual glassy phase, is visible at temperatures corresponding to the

second exothermic DSC peak (650 K), where the residual glass crystallizes into

intermetallic compounds. The values of Tx1 and Tx2 evaluated from the viscosity

measurements (623 and 670 K) are in good agreement with the data determined from the

constant-rate heating DSC scans (613 and 669 K) given the different instruments used for

the experiments. This promises that by utilizing the viscous flow of the supercooled liquid

(SCL) above the glass transition temperature, bulk amorphous samples can be produced by

consolidation of the MA powder.

In order to understand the mechanism underlying the first crystallization event,

Johnson-Mehl-Avrami (JMA) analysis was done by carrying isothermal DSC

measurements at different annealing temperatures at 583, 588 and 593 K. The

corresponding DSC curves are shown in Figure 3.4.

Figure 3.4 Isothermal DSC scans for the mechanically alloyed Al85Y8Ni5Co2 powder performed at different annealing temperatures.

All the DSC traces show a single exothermic peak with a symmetric bell shape

after a certain incubation period (t0) which decreases with increasing the annealing

50

temperature. A prepeak can be observed during the incubation period. This might be

related to the growth of pre-existed nuclei. The presence of the exothermic peak suggests a

crystallization mechanism consisting of nucleation from the supercooled liquid rather than

a simple grain growth of already present nuclei. In fact, as reported by Chen et al. [Chen

1988], the isothermal calorimetric signal for a nucleation and growth process is an

exothermic peak with a maximum at a non-zero time whereas a grain growth process is

characterized by a monotonically decreasing signal.

As already mentioned in section 1.2.3, the crystallized volume fraction is directly

proportional to the fractional area under the exothermic peak [Málek 2000, Venkataraman

2007b]. Figure 3.5(a) shows the typical sigmoidal curves for different annealing

temperatures derived from Figure 3.4, which represent the crystallized volume fraction (X)

as a function of the annealing time. The curves become steeper with increasing annealing

temperature, indicating that the transformation proceeds faster as the temperature is

increased.

Figure 3.5 (a) Crystallized volume fraction (X) versus time (t) for the isothermal annealing of the mechanically alloyed Al85Y8Ni5Co2 powder and (b) corresponding Avrami plots (0,15 < X < 0.85) for different annealing temperatures.

According to the Johnson-Mehl-Avrami equation (equation 1.11), JMA plots have

been obtained by plotting ln[ln (1/(1-X))] versus ln(t-t0) for the different annealing

temperatures. Figure 3.5(b) shows the JMA plots for the Al85Y8Ni5Co2 powder. The

transformation range under consideration is 15 – 85 vol.% of crystallized material (0.15 <

X < 0.85). The data do not lie on a straight line and, instead, describe curves with a slope

(i.e. the Avrami exponent n) that is continuously changing in the entire transformation

range considered. Such a variation of n does not allow the use of the JMA analysis for the

51

modeling of the crystallization behavior of the Al85Y8Ni5Co2 powder. However, it suggests

that different mechanism may operate during the devitrification process.

Although amorphization by mechanical alloying was successfully achieved for the

Al85Y8Ni5Co2 powder, two main disadvantages preclude the use of this alloy for the

consolidation into bulk specimens. The amorphization of the Al85Y8Ni5Co2 powder by MA

requires an extremely long milling time (~ 500 h), which makes this processing step not

satisfactory, slowing down the entire production process. In addition, due to such a high Al

content, the powder yield is very low, generally not exceeding 5%. This is due to the soft

aluminum, which sticks to the walls of the vials and on the surface of the balls during

milling. Also, an important effect related to the preferential sticking of the pure Al is the

possible change in the chemical composition of the final alloy [Zhang 2007].

Alternatively to mechanical alloying of elemental powder mixtures, Al85Y8Ni5Co2

glassy powders can be produced by controlled milling of melt-spun glassy ribbons.

Accordingly, in the next paragraph results concerning the production and characterization

of Al-based glassy powders by milling of Al85Y8Ni5Co2 ribbons is presented. The

pulverization of the melt-spun ribbons was achieved by using proper milling conditions, i.e.

interval-milling at low intensity, corresponding to a rather low kinetic energy, and

performed at cryogenic temperature in order to retain their glassy structure and to avoid

sticking of the material to the milling tools due to the high ductility of the ribbons.

3.2 Al85Y8Ni5Co2 glassy powders by milling of melt spun ribbons

The constant-rate heating DSC scan (20 K/min) of the as-spun Al85Y8Ni5Co2

ribbon is shown in Figure 3.6. The curve is characterized by a distinct glass transition (Tg)

before three exothermic heat flow events indicating the transformation from the solid-state

glass into the supercooled liquid, followed by the supercooled liquid (SCL) region (ΔTx =

Tx1 - Tg) which is about 22 K.

These three exothermic heat flow events occur due to crystallization during heating.

The onsets of Tg and of the crystallization events (Tx1, Tx2 and Tx3) for as-spun ribbon are

538, 560, 602 and 656 K, respectively. The enthalpies of crystallization of as-spun ribbon

related to the exothermic DSC peaks are ΔH1 = 32.1 J/g, ΔH2 = 36.4 J/g and ΔH3 = 47.7 J/g.

52

Figure 3.6 DSC scans (20 K/min ) for the as-spun Al85Y8Ni5Co2 ribbon and ribbon ball milled for 5 h.

The DSC scan of the Al85Y8Ni5Co2 ribbon is rather different with respect to the MA

powder with the same nominal composition (Figure 3.2). This is most likely due to a

change in the chemical composition in the MA process, resulting from the preferential

sticking of Al during milling.

Figure 3.7 XRD patterns (Cu-Kα) for the as-spun Al85Y8Ni5Co2 glassy ribbon and the ribbon heated at 20 K/min to completion of the first (580 K ) and second (630 K ) crystallization events.

The structural evolution during heating of as-spun ribbon is shown in Figure 3.7.

The XRD pattern of the as-spun ribbons shows the typical broad maxima characteristic for

amorphous materials together with a broad diffraction peak at about 2θ 44°, most likely ≈

53

due to the formation of a small amount of fcc Al during melt spinning. When the sample is

heated to 580 K, i.e. above the first crystallization peak, the XRD pattern displays the

formation of fcc Al (space group Fm3m) [Villars 1985]. Additionally, an overlapping

broad diffuse maximum due to the residual amorphous phase can be observed. XRD

pattern after heating up to the completion of the second exothermic DSC peak (630 K)

leads to formation of crystalline peaks the formation of the Al3Y phase (space group R-3m)

[Villars 1985].

The melt-spun ribbons were then ball-milled to produce glassy powders. The

pulverization of the ribbons was performed at cryogenic temperature (about 77 K) in order

to retain their glassy structure and to avoid sticking of the material to the milling tools

[Calin 2004]. Due to the low milling temperature, the melt-spun ribbons can be easily

pulverized. The yield of the powder obtained by this method is rather high, exceeding 90%.

Figure 3.8 XRD patterns (Cu-Kα radiation) for the as-spun Al85Y8Ni5Co2 glassy ribbon and ribbon ball milled for 5 h.

The effect of milling on the thermal stability and the microstructure of the melt-

spun Al85Y8Ni5Co2 ribbon are shown in Figure 3.6 and Figure 3.8, respectively. Milling for

5 h does not change the multi-step crystallization behavior characterizing the as-spun

ribbon. The values of Tg, Tx1, Tx2,, and Tx3 for the milled ribbon are 543, 562, 604 and

653 K and are, therefore, only slightly changed with respect to the as-spun ribbon. The

enthalpies of crystallization were found to be ΔH1 = 30.4 J/g, ΔH2 = 27.2 J/g and ΔH3 =

41.9 J/g.

54

This indicates that the first crystallization event is marginally affected, whereas the

subsequent events are much more influenced by the milling treatment. The mechanical

deformation does not induce crystallization of the glass, as illustrated by the XRD pattern

of the ribbon milled for 5 h (Figure 3.8), which, besides the broad diffraction peak already

observed in the as-spun ribbon, does not show additional crystalline precipitates. These

results indicate that Al85Y8Ni5Co2 glassy powders displaying strikingly similar structure

and crystallization behavior in comparison to the parent as-spun sample can be produced

by pulverization of glassy precursors by carefully controlling the milling conditions.

The thermal stability investigations of the milled powders reveal a distinct glass

transition followed by a supercooled liquid region. In this region the powders exhibit a

deformation regime characterized by a viscous flow behavior [Eckert 1997] that may allow

the production of bulk samples by hot consolidation at temperatures within the range of the

supercooled liquid region [Eckert 1997]. The viscosity of milled ribbons is shown in

Figure 3.9 evaluated by parallel plate rheometry according to the Stefan’s equation shown

in Figure 2.5.

Figure 3.9 Temperature dependence (heating rate 20 K/min) of the viscosity of the supercooled liquid for the single-phase Al85Y8Ni5Co2 glassy ribbon milled for 5 h.

The curve displays a decrease of viscosity with increasing temperature from about

1.4 x 1010 Pa s at about 400 K to 2.1 x 109 Pa s at 520 K, most likely due to structure

relaxation. As the glass transition temperature is reached and the glassy solid transforms

into the SCL (above 520 K) the viscosity displays a stronger decrease to the minimum

55

value of 2.5 x 108 Pa s at 565 K. At about 565 K the crystallization sets in and the viscosity

abruptly increases with increasing temperature, indicating the loss of liquid-like behavior.

The Al85Y8Ni5Co2 glassy powders from the milling of melt-spun ribbons were then

consolidated by hot pressing (HP) into cylindrical samples of 10 mm diameter and

5–10 mm length. The viscosity data together with the results from DSC were used to select

the proper consolidation parameters. Hot pressing was performed at 550 K using a pressure

of about 500 MPa. No extrusion of the hot-pressed specimens was possible due to the

insufficient viscous flow. However, the bulk HP samples are characterized by a density of

3.162 g/cm3, which is 96.7% of the density of the starting cylindrical rods of the pre-alloy.

The Al85Y8Ni5Co2 hot-pressed bulk specimens produced from milled ribbons are thus

characterized by a relatively high porosity, as shown by the SEM image in Figure 3.10.

Figure 3.10 SEM image of the hot-pressed single phase Al85Y8Ni5Co2 glassy ribbon.

In order to reach a higher density and, consequently, to improve the ductility of the

samples, the Al85Y8Ni5Co2 glassy powders from milled ribbons were blended with 50 and

70 vol.% of Al in order to produce glass-reinforced metal matrix composites (MMCs). To

obtain a homogeneous dispersion of the glass reinforcement, the blended powders were

milled for 10 minutes and then consolidated by hot pressing followed by hot extrusion.

Consolidation was performed at 520 K, in order to take advantage of the viscosity drop in

the supercooled liquid regime, using a pressure of 500 MPa. By using these consolidation

parameters, extrusion was performed in 10 minutes, a sufficiently short consolidation time

in order to avoid crystallization of Al-based glassy powders [Calin 2004]. For comparison

purposes, a bulk specimen was produced by extrusion of pure Al powders using the same

consolidation parameters as used for the MMCs.

56

Figure 3.11 DSC scans (20 K/min) for the as-spun Al85Y8Ni5Co2 ribbon, ribbon ball milled for 5 h and composite samples with 50 and 30 vol.% of glass reinforcement.

The DSC scans of the MMCs with 50 and 30 vol.% of glassy phase are shown in

Figure 3.11. The presence of pure Al does not change the overall crystallization behavior.

In fact, the values of Tg, Tx1, Tx2, and Tx3 were found to be 269, 286, 329 and 377 K for the

sample with 50 vol.% of glass and 272, 287, 330 and 378 K for the sample reinforced by

30 vol.% of glassy phase. These values are remarkably similar to the as-spun as well as to

the milled ribbons. The enthalpies of crystallization are ΔH1 = 13.6 J/g, ΔH2 = 11.5 J/g and

ΔH3 = 18.1 J/g for the sample with 50 vol.% glass and ΔH1 = 9.9 J/g, ΔH2 = 8.1 J/g and

ΔH3 = 12.7 J/g for the sample containing 30 vol.% of glassy phase, which, after

normalization by the vol.% of glass reinforcements, gives similar values with respect to the

single-phase milled ribbons.

As a typical example of the structure of the consolidated composites, Figure 3.12

shows the XRD pattern of the MMC with 50 vol.% of Al85Y8Ni5Co2 milled ribbon. The

pattern is characterized by few narrow diffraction peaks belonging to the fcc Al phase

together with the broad maximum belonging to the glassy phase at about 2θ ≈ 39°. This

indicates that no crystallization of the glass occurred during consolidation of the

composites.

57

Figure 3.12 XRD pattern (Cu-Kα radiation) for the hot-pressed and hot extruded composite with 50 vol. % of Al85Y8Ni5Co2 glass reinforcement

Figure 3.13 Relative density of the consolidated samples as a function of the volume of glass reinforcement.

Figure 3.13 shows the relative density of the consolidated samples as a function of

the volume fraction of glass reinforcement. The relative density of the specimens with

respect to the density of the starting materials used for the melt spinning experiments (i.e.,

the cylindrical rods of crystalline intermetallic compounds) decreases from 99.2% for the

58

sample with 30 vol.% of glass reinforcement to 97.4% for the specimen with 50 vol.% of

glassy phase and, finally, to 96.7% for the single-phase glass (100 vol.%).

A similar behavior was reported for Al-based MMCs reinforced with SiC particles

[Chung 1999, Slipenyuk 2006] and was ascribed to clustering of the reinforcing particles

[Slipenyuk 2006]. Figures. 3.14(a) and 3.14(b) show SEM images of the composites with

50 and 30 vol.% glass reinforcement, respectively. The images display a homogeneous

distribution of flake-shaped particles (the glassy phase) dispersed in the fcc Al matrix. No

porosity is visible, further corroborating the high density of the samples.

Figure 3.14 SEM micrographs for the hot extruded composites with (a) 50 vol.% and(b) 30 vol.% glass reinforcement.

On the other hand, the SEM image of the single-phase Al85Y8Ni5Co2 glassy

specimen produced by hot pressing of the milled ribbons (Figure 3.10) displays a large

number of pores. This indicates that incomplete bonding between the particles has

occurred during consolidation of the single-phase glass, leading to a rather poor

densification of the material.

Figure 3.15 shows a typical room temperature uni-axial compression true stress-

strain curves under quasistatic loading for the single-phase glass produced by HP of the

milled ribbons and for the composite materials along with the curve for the hot extruded

pure Al. The HP specimen of the single-phase glass exhibits an elastic regime of 1.45% up

to a stress of about 400 MPa but no appreciable ductility, most likely due to the residual

porosity of the sample that may initiate cracks leading to the early failure of the material.

The observed fracture strains exceed 40% for all the composite materials. However, due to

the strong softening characterizing the composite specimens after reaching the compressive

strength (the maximum compressive stress which the material is capable of sustaining

59

[ASTM 2004], the compression tests shown in Figure 3.15 were stopped after reaching the

maximum stress and before fracture occurrence.

Figure 3.15 Room temperature compression stress-strain curves (strain rate of 1x10-4 /sec) for the hot-pressed and hot extruded pure Al, hot pressed and hot extruded composite with 30 vol.% of Al85Y8Ni5Co2 glass reinforcement, hot-pressed and hot extruded composite with 50 vol.% glass reinforcement and hot-pressed single-phase Al85Y8Ni5Co2 glassy powder.

Pure Al exhibits an elastic regime of 0.2% before yielding, which occurs at about

75 MPa. After yielding the stress increases with increasing strain and the sample exhibits

work-hardening up to the compressive strength of 155 MPa, reaching a strain at maximum

stress of about 25%. The mechanical properties of pure Al are remarkably increased by the

addition of the glass reinforcement. The specimen containing 30 vol.% glass displays an

elastic regime of 0.3% and a yield strength of 120 MPa. The compressive strength is raised

to 255 MPa while retaining a strain at maximum stress of about 10%. When the amount of

glassy phase is further increased to 50 vol.% the elastic range is still 0.3% while the yield

and compressive strength are further increase to 130 and 295 MPa, respectively, and the

strain at maximum stress is found to be about 7%.

The prediction of the overall mechanical properties of a composite from the

properties of the single constituents is an important prerequisite for the material design and

application. Among the different methods for estimating the mechanical properties of a

composite, the rule of mixtures (ROM) is the simplest and most intuitive [Kim 2000, Kim

2001]. The ROM considers the properties of the composite as volume-weighted averages

of the components properties and assumes that the components are non-interacting during

60

deformation (Behavior of glass in the sintering of metal-glass materials). This approach

has been extensively used to model the mechanical properties of fiber-reinforced matrix

composites [Chawla 1998, Kelly 1972]. Two ROM methods have been widely employed

to predict the mechanical properties of composites [Kelly 1972, Kim 2000]: (i) the Voigt

model, based on the equal strain assumption, and (ii) the Reuss model, based on the equal

stress assumption. Although these models have been derived for the elastic properties of

composites, they have been also used for the overall plastic regime [Bruck 1999, Kim 2000,

Louzguine 2002b, Mileiko 1969].

The Voigt or iso-strain model assumes that the two components, matrix and

reinforcement, experience the same strain during deformation [Chawla 1998]. For the

stress of the glass-Al matrix composite [Kim 1999], this can be written as:

σc = Vgl·σgl + VAl·σAl , (3.1)

where V is the volume fraction, σ is the strength and the subscripts c, gl and Al indicate the

composite, the glass reinforcement and the Al matrix, respectively. It is often observed that

the strength of a composite is lower than predicted by the Voigt model [Sarkar 1982]. This

is generally attributed to (i) inadequate bonding, (ii) porosity and (iii) inherent material

defects, e.g. cracks [Sarkar 1982]. Therefore, the iso-strain treatment represents the upper

bound.

The lower bound is given by the Reuss or iso-stress model, which assumes that the

composite exhibits equal stress in the two components [Chawla 1998]. For the stress, this

can be written as:

σc = 1−

⎟⎟⎠

⎞⎜⎜⎝

⎛+

Al

Al

gl

gl VVσσ

. (3.2)

The effect of the porosity on the mechanical properties of a composite can be taken

into account by considering the volume of the composite material to be made up of three

different volumetric components, i.e. reinforcement, matrix and porosity [Madsen 2003]:

Vc = Vgl + VAl·+ Vp , (3.3)

where the subscript p denotes the porosity.

61

Figure 3.16 (a) Compressive strength and (b) strain at maximum stress (evaluated from Figure 3.15) as a function of volume percent of glass reinforcement for the samples: (■) hot extruded, (□) hot-pressed, (○)single-phase Al85Y8Ni5Co2 melt-spun glassy ribbon from reference [Louzguine 2002b], and (∇) calculated from Equation. 3.2.

The values of the maximum stress and the strain at maximum stress as a function of

the amount of the Al85Y8Ni5Co2 glass reinforcement is shown in Figures. 3.16(a) and

3.16(b) together with the values of the glassy ribbon [Louzguine 2002b] and of the single-

phase glass consolidated by hot pressing (present work). The strength of the samples

strongly deviates from the Voigt model (dotted line) and, instead, can be fitted well by

using the Reuss model (corrected for porosity by equation (3.3)) (dashed line). This

behavior indicates that the compressive strength obeys the iso-stress model. The significant

difference in strength observed between the single-phase Al85Y8Ni5Co2 glass consolidated

by hot pressing (400 MPa) and the melt-spun glassy ribbon with the same composition

(1250 MPa) [Louzguine 2002b] cannot be exclusively ascribed to the effect of porosity. In

fact, the maximum stress of a melt-spun ribbon with the density of the HP glass calculated

by equation (3.2) should exceed 1000 MPa (open triangle in Figure 3.16(a)). Most likely,

the considerably low strength of the HP sample is due to inadequate bonding between the

particles.

Figure 3.16(b) shows the strain at maximum stress for the different composite

materials. No values for the strain are available for a fully dense glassy specimen (such as

for the melt-spun glassy ribbon in Figure 3.16(a)). Therefore, the strain value of the low-

density single-phase Al85Y8Ni5Co2 glass consolidated by HP was used in the data fitting.

Similarly, to the maximum stress values in Figure 3.17(a), the corresponding strain at

62

maximum stress can be well fitted using the iso-stress Reuss’s model. Although lower than

predicted by the Reuss model, the strain of the HP specimen follows the iso-stress

treatment within the experimental error. This indicates that the poor particle bonding is less

significant in affecting the strain of the single-phase glass.

The validity of the Reuss model for the description of the composites studied in the

present work is justified by the following considerations. The milled ribbons used as

reinforcements are in the form of flake-shaped particles. Therefore, they can be treated as

short fibers as a first approximation. Although not aligned along the same direction, they

tend to lie on the same plane (see Figures. 3.16(a) and 3.16(b)), which is normal to the

applied stress during the compression test. Therefore, such composites can be considered

as a random fiber array in a matrix deformed perpendicularly to the fiber direction, the

latter being a requirement for the application of the iso-stress model [Chawla 1998]. In

addition, it has been reported that the Voigt model fits the data well for a high volume

fraction of reinforcement [Kim 2001]. In this case, the deformation affects the

reinforcements as a consequence of the smaller distance between the particles/fibers [Kim

2001]. On the other hand, the Reuss model works well for small volume fractions and

longer distances between the reinforcements, where the deformation mainly occurs in the

soft matrix [Kim 2001]. In the current work, the amount of reinforcements is relatively low

(≤ 50 vol.%) and the distance between the particles is large (> 50 μm). Therefore, most

likely the plastic deformation occurs mainly in the matrix and the mechanical properties

obey the iso-stress model.

63

Chapter 4: Synthesis and characterization of high-strength Al-

based alloys by consolidation of gas-atomized powders

The results shown in Chapter 4 clearly demonstrate that powder metallurgy

methods are suitable for the production of Al-based glassy and glass-composite materials

with high strength combined with considerable ductility. However, the production of Al-

based glassy powders by mechanical alloying of pure elements yields poor output, which

limits the use of MA for the production of Al-based powders with high Al contents. In

order to circumvent this limitation, Al-based amorphous powders can be produced by

pulverization of the melt-spun glassy ribbons. Although this technique allows the

production of larger quantities of high quality material (i.e. constant composition and

microstructure combined with low contamination levels) compared to MA, the procedure

is rather complex and requires the detailed characterization of every single melt-spun

ribbon. On the other hand, gas atomization offers the possibility to easily produce large

quantities of powders with homogeneous properties (e.g. structure and thermal stability)

along with a uniform size distribution of particles. Accordingly, in this chapter amorphous

and partially amorphous Al-based gas-atomized powders (GAP) with different

compositions have been used as precursors for the production of high-strength bulk

samples. The chapter is arranged in three sections. The first section deals with the detailed

study of Al84Gd6Ni7Co3 powders and consolidated samples with special emphasis given to

the crystallization behavior of the powders, their consolidation into bulk specimens and the

corresponding mechanical properties. The second section focuses on the consolidation and

mechanical behavior of Al87Ni8La5 produced by SPS. Here, a detailed study has been

carried out to understand the remarkable deformation behavior of the bulk samples. Finally,

the third part shows the characterization, synthesis and mechanical properties of the

Al90.4Y4.3Ni4.4Co0.9 alloy, which displays promising tensile properties comparable to

commercial high-strength Al-based alloys.

4.1 Gas-atomized Al84Gd6Ni7Co3 powder

Structural and thermal characterization. Figure 4.1.1 shows typical SEM

micrographs of the as-atomized Al84Gd6Ni7Co3 powder. The morphology of the powder is

smooth and round. Small-sized (< 1 µm) satellites are found around large particles (Figure

64

4.1.1(b)). Due to turbulent atomization conditions, both coarse and fine particles, arising

from efficient secondary breakup of particles during atomization. During in flight

collisions, the interaction results in welding of finer size particles to larger particles

[Özbilen 1999]. Smaller size powder particles with dimensions below 10 µm do not show

any feature, whereas larger particles display the presence of precipitates, suggesting that

partial crystallization occurred during gas atomization.

Figure 4.1.1 (a) and (b) SEM micrographs of the as-atomized Al84Gd6Ni7Co3 powder.

The XRD pattern of as-atomized Al84Gd6Ni7Co3 powder (Figure 4.1.2) exhibits a

broad maximum at 2θ angles between 35 and 55° which is typical for glassy materials

along with tiny Bragg peaks due to the presence of small amounts of fcc Al and

orthorhombic Al19Gd3Ni5 (space group Cmcm) [Gladyshevskii 1992] phases.

Figure 4.1.2 XRD patterns (Co Kα) for the gas-atomized Al84Gd6Ni7Co3 powder, in the as-atomized state and after heating to 633 and 873 K.

65

This might be due to the cooling rate during gas atomization (about 102 - 103 K/s)

[Suryanarayana 1991], which is not sufficiently high to suppress the formation of

crystalline phases, as already observed for gas-atomized Al-Ni-Mm (Mm = misch metal)

and Al-Ni-La powder [Ohtera 1992]. This is in agreement with the SEM micrographs in

Figure 4.1.1, which show the presence of precipitates for particles with larger size.

Figure 4.1.3 (a) Isochronal DSC scans taken at different heating rates and (b) corresponding Kissinger plots for the evaluation of the activation energies related to the first and second crystallization events.

Figure 4.1.3(a) presents the isochronal DSC scan for the Al84Gd6Ni7Co3 glassy

powder as a function of temperature taken at different heating rates (φ). The curves exhibit

a distinct glass transition with onset temperature Tg, followed by the supercooled liquid

region ΔTx before two crystallization events with onset and peak temperatures Tx and TP,

respectively, occur at higher temperature. The values TP at different heating rates are

summarized in Table 4.1.1.

Table 4.1.1. Summary of the results from the isochronal DSC experiments.

Heating Rate φ(K/min)

TP1 (K) TP2 (K)

10 578 ± 2 647 ± 2 20 588 ± 2 661 ± 2 40 593 ± 2 671 ± 3 60 595 ± 3 676 ± 3 80 599 ± 3 682 ±3

The peak temperatures TP1 and TP2 shift towards higher temperatures as the heating

rate is increased from 10 to 80 K/min. The supercooled liquid region ranges between 15

and 20 K, which is similar to what was observed for other Al-based amorphous systems,

such as Al-Y-Ni-Co [Inoue 1998].

66

The XRD pattern (Figure 4.1.2) of the sample annealed up to the completion of the

first exothermic DSC peak (633 K) reveals diffraction peaks from fcc Al and a broad peak

from the Al19Gd3Ni5 intermetallic phase together with the broad maxima of the residual

amorphous phase. The diffraction peaks of fcc Al are remarkably broad, which indicates

that this phase has nm-sized dimensions. A similar crystallization behavior has been

observed for different Al-based amorphous alloys, such as Al-Ni-RE and Al-Nd-Ni-Co

[Gao 2008, Huang 2008, Ye 2000]. When the sample is heated up to 873 K, i.e. far above

the second crystallization peak, the XRD pattern (Figure 4.1.2) shows the diffraction peaks

from fcc Al, Al19Gd3Ni5 and Al9Co2 (space group P21/a) [Douglas 1950] intermetallic

phases. No amorphous phase is visible at this stage, indicating that complete crystallization

of the glass occurs during the heat treatment.

Figure 4.1.4 Viscosity curve of as-atomized the Al84Gd6Ni7Co3 powder at a heating rate 10 K/min.

10 K/min) of the as-atomized powder. At the glass transition temperature (~550 K), where

the gla

scans taken at different heating rates using the

Kissing

Figure 4.1.4 shows the temperature dependence of the viscosity (heating rate

ssy solid transforms into the SCL, the curve displays a strong viscosity drop. At 590

K crystallization sets in and the viscosity abruptly increases with increasing temperature,

indicating the loss of liquid-like behavior.

The activation energy (Ea) for the crystallization process in Figure 4.1.3(a) was

evaluated from constant-rate heating DSC

er method (equation 1.10). By plotting ln(φ/T2P) versus (1/Tp), a straight line with

slope Ea / R is obtained (Figure 4.1.3(b)). The activation energies corresponding to the

67

crystallization events are Ea1 = 283 ± 2 kJ/mol and Ea2 = 210 ± 4 kJ/mol. These values are

comparable to the values attributed to self diffusion of aluminum [Volin 1968] suggesting

that the nucleation is diffusion controlled process.

In order to understand the mechanism underlying the first crystallization event,

isothermal DSC measurements were carried out at different annealing temperatures

ranging from 568 to 578 K and the corresponding curves are shown in Figure 4.1.5(a). All

curves show a single exothermic peak with an almost symmetric bell shape and an

incubation time (τ) that decreases with increasing annealing temperature. The inset in

Figure 4.1.5(a) shows the typical sigmoidal curves for different annealing temperatures

derived from time Figure 4.1.5(a), which represent the crystallized volume fraction as a

function of annealing time (see section 1.2.3). The curves become steeper with increasing

annealing temperature, indicating that the transformation proceeds faster as the annealing

temperature is increased.

Figure 4.1.5 (a) Isothermal DSC scans taken at different annealing temperatures and (inset) crystallized volume fraction (X) vs. time (t) and (b) Avrami plots (in the range 0.10 < X0.85) calculated from the isothermal DSC scans in Figure 4.1.4(a).

ing to the analysis

describ

<

Figure 4.1.5(b) shows the JMA plots for the Al84Gd6Ni7Co3 as-atomized powder in

the transformation range 10 – 85 vol.% (0.10 < X < 0.85) accord

ed in section 1.2.3. The Avrami exponent decreases from 3.0 ± 0.07 for isothermal

annealing at Ta = 568 K to 1.6 ± 0.04 for Ta = 578 K. Values of n of about 3.0 may be

related to a transformation mechanism characterized by diffusion controlled three

dimensional growth and increasing nucleation rate, whereas n = 1.6 suggests almost zero

nucleation rate [Avrami 1939, 1940, 1941]. This behavior can be understood by

considering the phase formation during crystallization. The first crystallization event is

68

characterized by the formation of nanocrystalline fcc Al (Figure 4.1.2), most likely through

a nucleation and growth mechanism, as suggested by the isothermal bell-shape peak in

Figure 4.1.5(a). Primary formation of nanocrystalline fcc Al from the amorphous phase is

characterized by an extremely high nanocrystal density of about 1021 – 1023 m-3 [Perepezko

2007], which may explain the high value of n = 3.0 observed for low annealing

temperatures. However, each fcc Al nanocrystal formed rejects the solute elements Gd, Ni

and Co into the residual amorphous matrix, thus reducing the driving force for the

formation of additional fcc Al and significantly reducing the nucleation rate [Allen 1998],

in accordance with the small value of the Avrami exponent (n = 1.6) observed for high

annealing temperatures.

Figure 4.1.6 Arrhenius plot for the isothermal activation energy of the Al84Gd6Ni7Co3 powder.

calculated from the intercept of the JMA plots (Figure 4.1.5(b)) (see equation 1.11). Under

isother

The reaction rate constant KT is a function of annealing temperature can be

mal conditions, the Arrhenius equation (equation 1.8) is often used to calculate the

activation energy for crystallization of an amorphous alloy [Scott 1977]. Figure 4.1.6

shows the plot of ln(KT) versus (1000/T), which yields a straight line whose slope gives the

activation energy of crystallization. The activation energy calculated by this method is

271 ± 3 kJ/mol. This value is remarkably similar to the activation energy calculated by the

Kissinger method (283 ± 2 kJ/mol).

69

Consolidation and mechanical properties. In order to study the influence of

temperature on densification and microstructural evolution, cylindrical samples (10 mm

diameter and 10 mm length) were produced from gas-atomized Al84Gd6Ni7Co3 powder by

hot pressing at 573, 623, 673 and 723 K with 3 minutes dwelling time. During hot pressing,

the change in the sample dimensions (shrinkage) occurs in the axial direction with

increasing the temperature. The corresponding lateral expansion does not occur due to the

constraints imposed by the wall of the die. The height variation of the powder bed in the z-

direction during hot pressing allows to understand the densification and the active sintering

temperatures [Garay 2010, German 1996]. For the present HP experiments, the

instantaneous variation of the powder bed (corrected for the dimensional changes of the

HP equipment) has been continuously measured. The shrinkage in the sample during HP at

constant pressure is given by Δh/ho, where ho is the initial height and Δh is the change in

height during the process. The shrinkage rate is the first derivative of the shrinkage

⎟⎟⎠

⎞⎜⎛ Δhd

rate versus temperature plots are useful in identifying the most active sintering temperature

[

⎜⎝ ohdt

, and since time and temperature are linearly related during constant rate heating,

this is derived from the shrinkage versus temperature plot. The shrinkage and shrinkage

German 1996].

Figure 4.1.7 Shrinkage data for the hot-pressed Al84Gd6Ni7Co3 samples

The shrinkage and shrinkage rate plots for hot pressing of the Al84Gd6Ni7Co3

powder are shown in Figure 4.1.7. The shrinkage increases sharply with increasing

temperature up to 573 K. For temperatures above 573 K, the slope decreases remarkably

70

and finally, above 650 K, the shrinkage curve displays a plateau, which indicates that full

densification is reached. In the first steps of hot pressing, the densification progresses

through particle rearrangement, particle bonding, necking and neck growth along with

plastic flow. As temperature increases, the shrinkage occurs due to mass transport by grain

boundary, surface and volume diffusion processes [German 1996]. The shrinkage rate

curve reveals that the maximum sintering rate is at about 573 K, which is within the super-

cooled liquid region. This implies that the shrinkage rate is maximum above the glass

transition region, where the sample experiences viscous flow behavior.

The density of the Al84Gd6Ni7Co3 samples hot-pressed at 573, 623, 673 and 723 K

is shown in Figure 4.1.7. The results reveal that the density is directly related to the

shrinkage curve and increases with increasing sintering temperature, reaching a plateau for

sintering at temperatures ≥673 K.

Figure 4.1.8 XRD patterns (Co-Kα radiation) of Al84Gd6Ni7Co3 powder: as-atomized, and

The XRD patterns of the HP samples are shown in Figure 4.1.8. The XRD pattern

of the

bulk sample consolidated by hot pressing at 573, 623, 673 and 723 K.

sample HP at 573 K reveals a broad maximum at 2θ ≈ 43° due to the residual

amorphous phase together with the broad diffraction peaks from fcc Al. Besides the

amorphous maximum and the diffraction peaks from fcc Al, less intense peaks from the

Al19Gd3Ni5 and Al9Co2 phases can be observed in the XRD pattern for the sample HP at

623 K. The XRD patterns for the samples hot-pressed at 673 and 723 K show the

diffraction peaks from fcc Al, Al19Gd3Ni5 and Al9Co2 phases. No broad maximum due to

the amorphous phase is visible at this stage, which indicates complete crystallization of the

71

glass. These diffraction peaks are rather broad, suggesting that the phases formed are of

nano or ultra-fine dimensions. The diffraction peaks are sharper for the samples HP at

723 K than the peaks observed in the samples HP at 673 K, which implies that larger grain

growth occurs at 723 K.

Figure 4.1.9 (a) OM and (b)-(d) SEM micrographs of sample hot-pressed at 573 K.

In order to analyze the effect of the densification on the microstructure of the

Al84Gd

e observed at the particle interfaces along with

particle deformation to an irregular “polygonal” shape due to the viscous flow behavior

6Ni7Co3 specimens, the microstructure of the bulk samples hot-pressed at different

temperatures have been investigated by OM, SEM and TEM and the corresponding

micrographs are shown in Figures 4.1.9 – 4.1.12. Figure 4.1.9 shows the OM and SEM

micrographs of the sample hot-pressed at 573 K, where the highest shrinkage rate occurs

(Figure 4.1.7). The OM micrograph (Figure 4.1.9(a)) reveals that densification occurs at

this sintering temperature. However, a large amount of porosity (≈2%) is clearly visible

(Figure 4.1.9(b)), which indicates that full densification is not achieved, corroborating the

shrinkage results shown in Figure 4.1.7.

Initiation of necking can also b

72

(Figure

sitional contrast

(Figure

compared to the material HP at 573 K. Figures 4.1.10(b) and 4.1.10(c) clearly

display

4.1.9(c)). In addition, a concentration gradient is visible as dark regions formed at

the particle boundaries near the pore regions (Figure 4.1.9(c)). The SEM-EDX results

reveal that these dark areas are rich in aluminum, which is in agreement with the XRD

pattern in Figure 4.1.8 showing the formation of Al for sintering at 573 K. Therefore, the

initial sintering stages discussed in section 1.3.3, i.e. initial bonding between the particles,

necking and concentration gradient, all occurs during sintering at 573 K.

The Al-rich concentration gradient (≈10% more than centre) mainly occurs around

the pores, whereas the necking areas do not show significant compo

4.1.9(d)). This implies that preferential diffusion of Al towards the pores takes

place during sintering. A possible explanation for this behavior might be related to the

evaporation-condensation transport mechanism. Vapor transport during sintering leads to

repositioning of atoms located on the particle surface. Evaporation occurs from a surface

and transport takes place across the pore space, leading to condensation on a nearby

surface. This causes mass transport towards the pore area [German 1996]. This type of

mass transport may be assisted by the viscous flow occurring during hot pressing. The

viscosity of the Al84Gd6Ni7Co3 powder drastically drops at temperatures above the glass

transition (Figure 4.1.4). The viscosity is inversely related to diffusivity [Cahn '96] and,

therefore, the low viscosity of the SCL at 573 K may enhance the diffusivity of Al leading

to an improvement of the mass transport towards the pore space. As a result, in the material

sintered at 573 K, viscous flow combined with evaporation-condensation of Al may lead to

the significant Al-rich concentration gradient around the pore space, as observed in the

Figures 4.1.9(c) and 4.1.9(d). However, these factors are not sufficient to obtain full

densification, as it is evident from the shrinkage curve for hot pressing at 573 K (Figure

4.1.7).

The SEM image of the sample HP at 623 K (Figures 4.1.10(a)) shows reduced

porosity

that necking at the particle interfaces is much more developed as compared to the

sample HP at 573 K (Figure 4.1.9). In addition, the powder particles maintained a

spherical/elliptical shape, which is in contrast to the “polygonal” shape characterizing the

sample HP at 573 K. This is most likely related to the different periods spent by the

samples at temperatures within the SCL region, where the material is characterized by a

viscous flow. While the sample HP at 573 K was kept isothermally within the SCL region

for 3 minutes, the sample HP at 623 K (which is above the minimum viscosity in Figure

73

4.1.4) passed through the SCL region without dwelling time, spending about 30 seconds at

temperatures within the SCL region.

Figure 4.1.10 SEM micrographs of the sample hot-pressed at 623 K.

es and in

the cen

(a)) shows few

small-s

The neck growth visible in Figure 4.1.10 results in particle shape chang

tre-to-center approach of the powder particles [German 1996]. In addition, neck

growth occurs by particle coalescence, as shown in Figure 4.1.10(c). Beside neck growth,

nc/UFG dendrites with size of 50 to 200 nm can also be observed inside the power

particles (Figure 4.1.10(d)). The formation of small dendrites is due to the high processing

temperature (623 K), which causes crystallization of the amorphous phase.

The OM micrograph of samples hot-pressed at 673 K (Figure 4.1.11

ized pores together with bright regions at the particle interface, revealing good

bonding of the particles. The black interface between the particles visible in the SEM

micrographs (Figures. 4.1.11(b)–4.1.11(d)), corresponding to the bright areas in Figure

4.1.11(a), is not due to porosity.

74

Figure 4.1.11 (a) OM and (b)-(d) SEM micrographs of sample hot-pressed at 673 K.

SEM-EDX analysis reveals that the particle interface consists of fcc Al. Pore

rounding occurs at this stage, as shown in Figure 4.1.11(d). In addition to the improved

densification, HP at 673 K induces extended crystallization of the glassy phase, as

demonstrated by the formation of several rod-like nano-sized features, which can be

identified as the Al19Gd3Ni5 and Al9Co2 intermetallic phases observed by XRD (Figure

4.1.8).

The OM image of the sample hot-pressed at 723 K (Figure 4.1.12(a)) shows no

visible porosity. The SEM micrographs (Figures 4.1.12(b) and 4.1.12(c)) reveal that the

particles consist of agglomerates of bright rod-like particles and black regions of nm-scaled

dimensions. The TEM-BF images in Figure 4.1.12(d) shows the triangular area between

particles. TEM-EDX analysis reveals that the triangular area consists of UFG Al of about

200 to 300 nm size. The regions surrounding the triangular area consist of nano-sized Al

grains (black particles) and rod-like nano-sized features rich in Gd, Ni and Co of about

50 nm thickness and 200 nm length, most likely corresponding to the Al19Gd3Ni5 and

Al9Co2 compounds observed by XRD (Figure 4.1.8). This is corroborated by the work of

75

Gao et al. [Gao 2003] who observed a rod-like morphology for the ternary Al19Gd3Ni5

phase.

Figure 4.1.12 (a) OM, (b) and (c) SEM, (d) TEM-BF micrographs of sample hot-pressed at 723 K.

The hardness of the samples HP at 573, 623, 673 and 723 K reveal Vickers

microhardness of about 289 ± 15, 542 ± 10, 535 ± 8, 432 ± 10 HV, respectively. The

lowest hardness values observed in the samples HP at 573 K can be ascribed to the poor

bonding and presence of large pores characterizing this sample. On the other hand, for the

samples HP at higher temperatures, which display good bonding between the particles and

small residual porosity, the hardness values are rather high and show a decreasing trend

with increasing HP temperature. This may be related to the crystallization of the hard

amorphous phase and the formation of softer phases (e.g. Al). However, the hardness of

the present consolidated samples is much higher than for conventional Al-based alloys

(60 – 250 HV [Davis 1993]) due to their nanocrystalline/ultra fine-grained microstructure.

76

Figure. 4.1.13 Room temperature compression true stress-true strain curves for the Al84Gd6Ni7Co3 gas-atomized powder hot-pressed at different temperatures.

Typical room temperature uni-axial compressive true stress-strain curves for the HP

samples are shown in Figure 4.1.13. The samples HP at 623 K fractured in a brittle manner

at about 700 MPa without any yield or plastic strain. This early fracture might be due to the

presence of large pores and inadequate particle bonding, as observed in Figure 4.1.10. The

sample HP at 673 K exhibits a yield stress (0.2% offset) of about 1250 ± 10 MPa followed

by strain hardening up to the maximum stress of 1560 ± 5 MPa where fracture occurs at

3.5 ± 0.2%. The sample HP at 723 K exhibits lower yield and compressive strengths

(1150 ± 10 MPa and 1440 ± 7 MPa, respectively) with respect to the sample HP at 673 K,

however, the fracture strain slightly increases to 4 ± 0.4%. The stress-strain curves for the

samples HP at 673 and 723 K exhibit strain hardening behavior. The decreased values of

yield and compressive strengths for the samples HP at 723 K as compared to the samples

HP at 673 K may be linked to the larger grain size occurring at 723 K, as shown by the

XRD patterns in Figure 4.1.8. This may also explain the larger plastic deformation shown

by the sample HP at 723 K.

The strength levels of the samples HP at 673 and 723 K are three times larger than

for conventional Al-based high strength alloys [Davis 1993]. Such high strength levels are

most likely due to the multi-phase microstructure consisting of a uniform distribution of

high-strength nanocrystalline rod-shaped intermetallic phases and nanocrystalline Al

particles along with areas of deformable ultra fine-grained Al at the triangular areas.

77

Figure 4.1.14 SEM micrographs of fractured samples after compression testing of samples HP at 723 K.

Plastic deformation during compression testing of ultra fine-grained Al at the

triangular areas is corroborated by the fracture surface of the sample HP at 723 K (Figure

4.1.14(a)), revealing the formation of dimples at the triangular areas, indicative of ductile

fracture. On the otherhand, the areas with high density of intermetallic compounds display

brittle fracture characterized by intra-granular rupture as well as decohesion of the particles.

Although the Al regions at the triangular areas do not form a continuous network

throughout the specimen, they may nevertheless allow significant movement of

dislocations. This, together with the obstacle to the dislocation movement due to the

intermetallic phases, may explain the high strength levels combined with a limited but

distinct good plastic deformability of the present HP samples.

78

4.2 Gas-atomized Al87Ni8La5 powder

The results presented in the previous section indicate that consolidation of the

Al84Gd6Ni7Co3 powder into highly-dense bulk samples cannot be achieved without

extended crystallization of the material. Nevertheless, crystallization during consolidation

is not detrimental and leads to bulk samples with a remarkably high strength of about

1550 MPa. However, such a high strength is accompanied by a relatively limited room

temperature plastic deformation of about 3.5 to 4%. To further test the effectiveness of

powder consolidation as a method for the production of Al-based materials with enhanced

mechanical properties through the combined crystallization and consolidation of glassy

precursors, in this section Al87Ni8La5 gas-atomized powders are consolidated by SPS

above their crystallization temperature and their mechanical properties are investigated in

detail.

Figure 4.2.1 XRD pattern (Co-Ka radiation) of Al87Ni8La5 powders: as-atomized, after heating (heating rate 40 K/min) to different temperatures in the DSC, and bulk sample consolidated by SPS at 713 K.

The structure of the as-atomized powder, investigated by XRD, is presented in

Figure 4.2.1. The pattern displays the broad diffuse maxima, characteristic of amorphous

materials, together with few broad diffraction peaks corresponding to fcc Al and Al11La3

(space group - Immm) phases. This implies that the structure of the as-atomized Al87Ni8La5

powder is not completely amorphous, as already observed for the Al84Gd6Ni7Co3 powder

79

(section 4.1, Figure 4.1.2), and further confirms the limited glass-forming ability of Al-

based alloys with high Al content.

Figure 4.2.2 shows the constant-rate heating DSC scan (20 K/min) of the as-

atomized powder. The DSC curve exhibits two exothermic events due to crystallization

with onset temperatures Tx1 = 445 K and Tx2 = 612 K. This type of thermal behavior can be

observed in many Al-based glassy alloys with high Al contents above about 85 at.%, which

crystallize through two stages upon heating to elevated temperatures [Inoue 2001b]. The

enthalpies of crystallization related to the first and second exothermic DSC peaks are ΔH1

= 19 ± 3 J/g and ΔH2 = 77 ± 5 J/g, respectively. These values are very similar to those

reported for Al-Ni-La alloys with similar composition (22 and 80 J/g, respectively) [Ye

2000], suggesting that only a small fraction of material crystallized during gas atomization.

Figure 4.2.2 DSC scan and (inset) temperature dependence of the viscosity for the as-atomized Al87Ni8La5 powder.

Figure 4.2.1 shows the XRD patterns of the Al87Ni8La5 powder obtained after

continuous heating at 20 K/min to 573, 673 and 773 K. When the sample is heated up to

573 K, i.e. above the first crystallization peak, the XRD pattern displays the presence of

fcc Al and the simultaneous decrease of the diffuse amorphous maximum. Heating above

the second exothermic DSC peak (673 and 773 K) leads to the decomposition of the

residual glassy phase into the intermetallic compounds Al11La3 and Al3Ni.

Similarly to the MA Al85Y8Ni5Co2 (section 3.1), no clear glass transition prior to

crystallization can be observed in the DSC scan in Figure 4.2.2. In order to verify the

occurrence of the glass transition, the influence of temperature on the viscosity of the as-

80

atomized powder was studied by parallel plate rheometry (inset in Figure 4.2.2). The

viscosity decreases with increasing temperature from about 3 x 108 Pa s at 400 K to

5 x 106 Pa s at 470 K, which indicates the occurrence of the glass transition and the

transformation of the glassy solid into the supercooled liquid (SCL). At about 470 K

crystallization sets in and the viscosity abruptly increases with increasing temperature,

indicating the loss of liquid-like behavior.A second drop of viscosity is visible at

temperatures corresponding to the second exothermic DSC peak (610 K), where the

residual glass crystallizes into intermetallic compounds (Figure 4.2.1). It has been reported

for Al-based glasses that Tg and Tx increase significantly with increasing solute

concentration [Inoue 1998]. Although for the present Al87Ni8La5 powder the composition

of the residual glassy phase after the first crystallization event is not known, it is most

likely depleted in Al due to the primary formation of fcc Al, as shown in Figure 4.2.1. It is

thus plausible that the residual Al-poor glassy phase undergoes the glass transition at

higher temperature, explaining the second drop of viscosity.

The occurrence of two viscosity drops may be a considerable advantage for the

consolidation of the gas-atomized powder in assisting interparticle bonding and

densification. In order to take benefit of the double viscous flow behavior, during

consolidation by SPS a constant pressure of 500 MPa was applied from room temperature

through the crystallization events to the final sintering temperature (713 K), which was

held for about 5 minutes. The consolidation of the Al87Ni8La5 powder at high temperatures

can thus be considered as a combined in-situ devitrification and densification of the

powders.

The XRD pattern of the Al87Ni8La5 powder sintered at 713 K is also shown in

Figure 4.2.1. The structure consists of fcc Al together with two intermetallic compounds,

i.e. Al11La3 and Al3Ni. No trace of the amorphous phase is visible, indicating that complete

crystallization occurred during sintering. The diffraction peaks in Figure 4.2.1 are rather

broad, indicating that the phases formed are of nanoscale or ultra-fine dimensions. Indeed,

Rietveld structure refinement reveals an average grain size for the different phases ranging

between 100 and 200 nm. This is in agreement with the results reported for the

devitrification of melt-spun Al87Ni8La5 glassy ribbons [Sahoo 2005], which show the

formation of similar intermetallic phases with an average grain size of less than 140 nm.

81

Figure 4.2.3 (a) OM, (b) SEM, (c)-(e) TEM micrographs, and (f) EDX elemental mapping of the Al87Ni8La5 powder consolidated by SPS at 713 K.

The microstructure of the consolidated bulk material was studied by OM, SEM and

TEM, and the corresponding micrographs are shown in Figure 4.2.3. Optical microscopy

investigations (Figure 4.2.3(a)) reveal the formation of a bright interface layer between the

particles, which indicates that particle bonding has taken place during consolidation. The

SEM micrograph in Figure 4.2.3(b) shows that the particles mostly retain their original

spherical shape with a neck geometry characterized by center approach and particle

penetration [Exner 1979], as shown in the inset in Figure 4.2.3(b). TEM investigations of

the sintered sample (Figures. 4.2.3(c)–4.2.3(e)) show that the microstructure of the

particles consists of black areas, often continuously connected, with dimensions of about

200 – 300 nm along with bright areas with dimensions in the range 100 – 200 nm. The

TEM micrograph in Figure 4.2.3(e) and the corresponding EDX elemental mapping in

Figure 4.2.3(f) reveal that the black areas consist of fcc Al whereas the bright areas

82

comprise two phases: a Ni-rich phase with size below 100 nm and a La-rich phase with

dimensions of about 200 nm, which most likely correspond to the intermetallics Al3Ni and

Al11La3 observed by XRD (Figure 4.2.1). The black interface between the particles

(triangular areas) visible in the SEM and TEM micrographs (Figures. 4.2.3(b)–4.2.3(d)),

corresponding to the bright areas in Figure 4.2.3(a), is not due to porosity. TEM and EDX

analysis reveals that the particle interface consists of an fcc Al matrix along with several

bright particles with dimensions below 50 nm, most likely Al11La3 and/or Al3Ni phases.

A similar microstructure (i.e. Al at the triangular areas) has been observed for the

HP Al84Gd6Ni7Co3 sample (Figure 4.1.12). However, for the Al87Ni8La5 powder

consolidated by SPS the fcc Al phase at the triangular areas is continuously connected to

the particles through fcc Al channels (see arrows in Figure 4.2.3(d)) with thickness of

about 200 nm. This microstructure promises improved plastic deformation with respect to

the HP Al84Gd6Ni7Co3 sample as a result of the enhanced dislocation activity. Only few

pores are visible (indicated by arrows in Figure 4.2.3(c)), corroborating the high density of

the sintered samples evaluated by density measurements (~98 %).

Figure 4.2.4 (a) Compression true stress-true strain curves for the Al87Ni8La5 bulk material consolidated by SPS at 713 K (present work) and Al85Ni10La5 consolidated by SPS at 753 K [Sasaki 2008]. (b) Compression true stress-true strain curves for the Al87Ni8La5 material and pure Al: experimental data (lines) and values calculated by using the effective medium approach (points).

A typical room temperature uni-axial compression stress-strain curve for the

sintered sample is shown in Figure 4.2.4(a). The material exhibits an elastic regime of

0.7% and a yield strength (0.2% offset) of about 740 MPa followed by a region with strain

hardening up to the maximum stress of 930 MPa. After reaching the maximum, the stress

gradually decreases with increasing strain to about 640 MPa and fracture occurs at 27%

strain. Similar features have been recently reported for nanocrystalline Al-5 at.% Fe

83

consolidated by SPS [Sasaki 2007]. The sintered Al-Fe exhibits a compressive strength of

1045 MPa followed by a pronounced work softening-like behavior and a plastic strain of

about 30% [Sasaki 2007]. Plastic deformation of the sintered Al87Ni8La5 sample does not

lead to further densification, as demonstrated by the density of the specimen after

compression test that is reduced by about 1% with respect to the as-sintered specimen

(98% of the theoretical density).

Spark plasma sintering of the Al87Ni8La5 powder leads to a highly dense specimens

displaying high strength combined with remarkable plastic deformation. Such a behavior is

presumably due to the multi-phase microstructure consisting of soft fcc Al and high-

strength intermetallic compounds [Ohtera 1992]. However, the observed room temperature

plastic deformation is in contrast to what was reported for other Al-based alloys produced

by consolidation of gas-atomized powders [Inoue 2001a, Kawamura 2001], which display

a similar microstructure. For example, although fully crystallized Al-Ni-Y-Co samples

exhibit an extremely high compressive strength of 1420 MPa [Inoue 2001a, Kawamura

2001], the plastic strain is only about 1% [Inoue 2001a]. This is similar to recent results of

Sasaki et al. [Sasaki 2008] on fully crystallized nanocrystalline Al85Ni10La5 samples

produced by SPS of gas-atomized amorphous powders, which show a compressive strength

exceeding 1200 MPa, but no plastic deformation (Figure 4.2.4(a)).

Figure 4.2.5 SEM micrographs of a consolidated Al87Ni8La5 specimen: (a) as-sintered and (b) polished fracture surface after the compression test.

A possible explanation for the larger plastic deformation of the present Al87Ni8La5

with respect to the Al85Ni10La5 sample of Sasaki et al. [Sasaki 2008] related to the different

microstructures of the sintered samples. The sample of Sasaki et al. [Sasaki 2008] sintered

at 753 K (above the crystallization temperature) displays a microstructure consisting of fcc

Al regions surrounded by a large amount of intermetallic particles [Sasaki 2008]. The

regions of fcc Al appear to be confined and constrained by the intermetallic phases, which

84

presumably leads to the observed lack of deformation capability as a result of the limited

dislocation nucleation and movement. In contrast, in the present sample the fcc Al regions

at the particle interface (inter-particle) are continuously connected to the particles through

intra-particle fcc Al channels (Figure 4.2.3), giving rise to a network of ultra fine-grained

(UFG) Al reinforced with nm-scale intermetallic particles, which extends over the entire

specimen. Within this structure, the fcc Al regions are not confined and, as a result, the

continuous network of fcc Al may allow the movement of dislocations, explaining the

remarkable plastic deformation. This is corroborated by the comparison between the

particle morphology in the sintered samples before and after compression (Figure 4.2.5).

Before testing, the particles have a regular spherical shape (Figure 4.2.5(a)), whereas after

compression they collectively assume a squeezed elliptical shape (Figure 4.2.5(b)) with

major axis perpendicular to the compression direction (indicated by arrows in Figure

4.2.5(b)). Assuming that the average local deformation (εp) of the particles is given by εp =

(l - r)/r,, where l is the major axis of the ellipse after compression and r is the radius of the

original undeformed spherical particle, the value of εp is about 17%. This deformation can

be ascribed to the intra-particle Al regions because intermetallic phases are typically brittle

at room temperature [Koch 1998] and, therefore, are not able to deform plastically.

Similarly to the surrounding particles, the inter-particle Al regions at the interface also

deform plastically in order to keep the geometrical integrity (Figure 4.2.5(b)).

The uniaxial strain-stress curve of deformable metallic materials can be expressed

by the Ramberg-Osgood equation [Clyne 1993]

n

y

y

EE

/1

⎟⎟⎠

⎞⎜⎜⎝

⎛+=

σσσ

ασε , (4.2.1)

where E is the Young’s modulus of the metallic material, n is the strain hardening

exponent, α a dimensionless constant (α = 3/7 is usually taken for Al-based alloys

[Wilkinson 2001], and the yield strength. Based on Equation (eq. 4.2.1), the strain-

stress curve can be successfully modeled by using the effective medium approach (EMA)

[

yσ

Kouzeli 2002, Nan 1996, Scudino 2009]. For the present Al-Ni-La sample, calculations

using Equation (equation 4.2.1) and EMA are performed to determine n of the fcc Al

(Figure 4.2.4(b)). The calculations are in good agreement with the experimental results in

the strain hardening part of the curve before the critical strain of structural instability,

giving a value of n of 0.16. On the other hand, the value of n for pure aluminum evaluated

85

from the strain-stress curve in Figure 4.2.4(b) is only 0.04. This indicates that the

Al87Ni8La5 sample is characterized by a more intense dislocation storage and interaction

compared to pure aluminum [Hull 2001]. Two main reasons may be responsible for this

behavior: (i) extensive generation of dislocation and (ii) dislocation movement limited by

constraint effects [HullBacon 2001]. When a crystal is plastically deformed, dislocations

are generated, moved and stored. Dislocation storage, which causes the material to work-

harden, occurs by mutual trapping or by accommodating the deformation incompatibility

between various parts in the deformed materials. The dislocations that are mutually trapped

are referred to as statistically stored dislocations [Ashby 1971] and their density Sρ is

difficult to estimate. The dislocations that are stored due to deformation incompatibility are

called geometrically necessary dislocations, Gρ . Gρ is mainly related to the thermal

mismatch strains imposed upon cooling down from the processing temperature ( ) and

to the strain gradient present during deformation ( ), i.e., [

thGρ

sgGρ sg

GthGG ρρρ += HullBacon

2001]. In the case of the present sample consisting of fcc Al and rigid intermetallics,

and can be expressed respectively as [

thGρ

sgGρ Arsenault 1986, Nix 1998]

⎟⎟⎠

⎞⎜⎜⎝

⎛++

−

Δ×Δ=

321

111)1(

4DDDbV

TCTEV

p

pthGρ , (4.2.2)

λερ

bsgG =

4 , (4.2.3)

where is the difference in the coefficient of thermal expansion between fcc Al and

the intermetallics,

CTEΔ

TΔ is the temperature differential upon cooling (about 700 K), is the

volume fraction of intermetallic phase (about 0.60), b is the magnitude of the Burger’s

vector (0.28 nm [

pV

Rooy 1979]), , and are the three-dimensional sizes of the

intermetallic phase (about 1 μm × 1 μm × 0.5 μm), ε is the strain and λ is the local length

scale of the deformation field or the size of fcc Al (about 500 nm, see Figure 4.2.3(c)). By

using of about 5 × 10-6 K-1, is estimated as 3.0 × 1014 m-2. From Equation (4.2.2)

follows that increases with the strain. At the applied strain of 2%, can reach a

value of about 5.7 × 1014 m-2. This indicates that a high density of dislocations nucleate in

the fcc Al during deformation, explaining why the present sample has a strain hardening

exponent much larger than that of pure aluminum.

1D 2D 3D

CTEΔ thGρ

sgGρ sg

Gρ

86

Figure 4.2.6 SEM images of a consolidated Al87Ni8La5 specimen: (a) as-sintered, (b) after 4% and 10% plastic deformation.

Although the Al network is able to accommodate a large amount of plastic

deformation, deformation of fcc Al through dislocation activity is nevertheless limited by

the presence of the rigid intermetallics, which most likely makes further deformation

increasingly difficult. The storage of dislocations then may cause stress concentration in

the constrained Al, resulting in the formation of microcracks between the particles. In

order to clarify this aspect, the microstructure of the Al87Ni8La5 consolidated samples at

different stages of plastic deformation (4 and 10%) was analyzed using SEM and the

corresponding images are shown in Figures. 4.2.6(b) and 4.2.6(c), respectively, together

with the microstructure of the as-sintered sample (Figure 4.2.6(a)). The SEM micrographs

in Figure 4.2.6 reveal interesting features. The contraction of the powder particles parallel

to the compression direction (indicated by arrows in Figure 4.2.6), already observed in

Figure 4.2.5(a), starts to be visible at a plastic deformation of 4% and becomes evident for

the sample deformed up to 10% (see for example the particles labeled 1, 2 and 3). At the

same time, microcracks are formed at the interface between the particles (particles labeled

4 and 5).

87

Only few cracks are clearly visible in Figure 4.2.6 most likely because the SEM

investigation was carried out on the sample surface and, therefore, the inter-particle Al is

only partially constrained. In the inner part of the sample a larger number of microcracks

may form even at lower strains due to the giant hydrostatic stress and the three dimensional

constraints [Goods 1979]. Nevertheless, Figure 4.2.6 clearly shows that microcracking

occurs during deformation. As a consequence of crack formation and resulting stress

relaxation, softening-like behavior can be observed in the strain-stress curve (Figure 4.2.4

(a)). For large deformations, microcracks readily coalesce to form a main crack that rapidly

propagates through the sample, finally leading to fracture. Figure 4.2.7 shows the fracture

surface of the compressed sample, where inter-particle fracture can be clearly observed.

Besides particle deformation and microcracking, the sample deformed up to 10%

displays profuse shear banding, as shown by the OM and SEM images in Figure 4.2.8.

Most likely shear banding accommodates the strain when the dislocation movement is

limited. The shear bands form an angle of about 45° with the compression direction

(indicated by arrows in Figure 4.2.8). Multiple shear bands are often observed during the

deformation of metallic glass composites [Hays 2000, He 2003]. The difference between

the present material and metallic

Figure 4.2.7 SEM micrograph of the fracture morphology after compression test for the consolidated Al87Ni8La5 specimen.

glass composites is that shear banding is typically the only deformation mechanism of

glassy composites, while dislocation motion is also responsible for the plastic deformation

of the present composite, as shown in Figure 4.2.8(d), where, within the shear bands, both

particles and inter-particle matrix are plastically deformed. The activation of multiple shear

bands together with the dislocation-associated deformation results in a composite having a

fracture strain of about 27%.

88

Figure 4.2.8 (a)-(b) OM and (c)-(d) SEM images of the consolidated Al87Ni8La5 sample plastically deformed up to 10%, displaying profuse shear banding.

89

4.3 Gas-atomized Al90.4Y4.4Ni4.3Co0.9 powder

As shown in the previous two sections, the consolidation of Al-based gas-atomized

powders yields highly-dense bulk samples characterized by high strength levels. The HP

Al84Gd6Ni7Co3 powder exhibits remarkably high strength of about 1550 MPa as compared

to about 930 MPa for the Al87Ni8La5 sample consolidated by SPS. However, the plastic

strain is significantly improved for the Al87Ni8La5 material. This is related to the higher Al

content of the Al87Ni8La5 alloy with respect to the Al84Gd6Ni7Co3 powder and to the

corresponding microstructural features: increasing Al content gives rise to enhanced plastic

strain along with a decrease of strength. To further investigate the effect of Al content on

the mechanical behavior of the consolidated samples, in this section, Al90.4Y4.4Ni4.3Co0.9

gas-atomized powders are consolidated by hot pressing and their mechanical properties are

investigated in detail. Additionally, samples were consolidated by hot pressing followed by

hot extrusion in order to test the tensile mechanical properties of the bulk sample and to

compare the tensile mechanical behavior with the results obtained in compression.

Figure 4.3.1 XRD pattern (Co-Kα radiation) of Al90.4Y4.4Ni4.3Co0.9 powder: as-atomized, after heating to 640 and 873 K in the DSC, bulk sample consolidated by hot pressing at 673 and 723 K and samples hot pressed and hot extruded at 723 K.

The microstructure of the as-atomized Al90.4Y4.4Ni4.3Co0.9 powder investigated by

XRD is shown in Figure 4.3.1. The pattern exhibits broad maxima typical for an

amorphous phase together with broad diffraction peaks of fcc Al. This indicates that the as-

atomized powder is a mixture of an amorphous phase and nanocrystalline Al, which is

90

similar to the structure observed for the Al87Ni8La5 gas-atomized powder (Figure 4.2.1).

The DSC scan (40 K/min) of the as-atomized Al90.4Y4.4Ni4.3Co0.9 powder is shown in

Figure 4.3.2. The DSC curve exhibits two exothermic peaks due to the crystallization of

the glass without any clear glass transition region. The onset temperatures of first and

second exothermic peaks are Tx1 = 600 K and Tx2 = 650 K, respectively.

Figure 4.3.2 DSC scan (40 K/min) of as-atomized Al90.4Y4.4Ni4.3Co0.9 powder.

The XRD patterns of the as-atomized powder after heating to different

temperatures are shown in Figure 4.3.1. The XRD pattern after heating to 640 K (above the

first crystallization event) shows broad diffraction peaks belonging to Al along with

diffraction peaks from the Al3Y and Al4Ni3 intermetallic compounds. In addition, an

extremely weak and broad maximum due to the residual amorphous phase can be observed.

The XRD pattern of the powder heated to 873 K (far above the second crystallization

event) shows the peaks from Al and intermetallic phases Al3Y and Al4Ni3 and Al9Co2. No

amorphous phase is visible at this stage, indicating that complete crystallization of the

glass occurs during the heat treatment at 873 K.

According to the results presented in section 4.2, excellent mechanical properties

can be obtained for the Al84Gd6Ni7Co3 powder by HP at 673 and 723 K. Therefore, the

same sintering temperatures (673 and 723 K) have been used to consolidate the

Al90.4Y4.4Ni4.3Co0.9 powder by HP. The XRD pattern of the Al90.4Y4.4Ni4.3Co0.9 sample hot-

pressed at 673 K (Figure 4.3.1) reveals the formation of the phases already observed in the

sample heated to 640 K in the DSC, i.e. Al, Al3Y and Al4Ni3 intermetallic compounds. As

well, the week broad maximum due to the presence of a residual amorphous phase can be

91

observed. The presence of residual amorphous phase is due to the sintering temperature

(673 K), which is below the second crystallization event. The XRD pattern of the sample

HP at 723 K (above the second crystallization peak in Figure 4.3.2) reveals the formation

of Al, Al4Ni3, Al3Y and Al9Co2 phases without any visible residual amorphous phase

(Figure 4.3.1).

Figure 4.3.3 (a)-(b) SEM micrographs (c)-(d) bright-field TEM micrographs of the Al90.4Y4.4Ni4.3Co0.9 samples hot-pressed at 673 K

Figures 4.3.3(a) and 4.3.3(b) show the SEM micrographs of the sample hot-pressed

at 673 K. The micrographs reveal the formation of a bright interface between the particles

along with black areas at the triangular areas, as already observed for the Al84Gd6Ni7Co3

and Al87Ni8La5 consolidated samples (Figures 4.1.12 and 4.2.3). The particles are not

homogeneous and can be considered as a composite microstructure consisting of grains of

about 1-3 μm size separated by a matrix with rod-like morphology (Figure 4.3.3(b)). EDX

analysis indicates that the bright interfaces are rich in Y, Ni and Co (≈ 40±10% rich than

the original atomic percent), which may correspond to the intermetallic compounds

92

observed by XRD (Figure 4.3.1). A similar composition has been observed for the rod-like

matrix between the particles. On the other hand, the 1-3 μm sized grains in Figure 4.3.3(b)

show a higher Al content (of about 5% of that of nominal Al content) that, considering the

XRD results in Figure 4.3.1, suggests a mixed microstructure consisting of Al and the

residual amorphous phase.

Figures 4.3.3(c) and 4.3.3(d) show the TEM-bright field micrographs

corresponding to the black regions between the particles (triangular areas) observed in

Figure 4.3.3(a). The figures reveal that the triangular areas are made of Al grains of about

100 – 300 nm and are surrounded by the rod-like features (the intermetallic compounds)

with 40 nm thickness and 300 to 500 nm length.

Figure 4.3.4 shows the SEM micrographs of the sample hot-pressed at 723 K.

Similarly to the sample HP at 673 K (Figure 4.3.3), the micrographs show the formation of

bright interfaces between the particles, indicative of good particle bonding, along with the

presence of the rod-like structures. The amount of rod-like structure in the sample hot-

pressed at 723 K is higher (≈ 30 ± 10 % from the image analysis) as compared to the

sample HP at 673 K. This is accompanied by the disappearance of the 1-3 μm sized grains

visible in Figure 4.3.3(b) and to the formation of a uniform microstructure throughout the

particle consisting of rod-like intermetallics and ultra fine-grained Al grains. The formation

of such a microstructure can be related to the high sintering temperature (723 K) which

leads to the full crystallization of the amorphous phase and to the formation of additional

fcc Al and intermetallic compounds from the residual amorphous phase (Figure 4.3.1).

Figure 4.3.4 SEM micrographs of hot-pressed Al90.4Y4.4Ni4.3Co0.9 samples at 723 K

Figure 4.3.5 shows the room temperature compressive stress-strain curves of the

samples consolidated by hot pressing at 673 and 723 K. The sample HP at 673 K exhibits a

93

yield and compressive strength of about 880 ± 10 MPa and 925 ± 2 MPa, respectively.

With increasing stress, the curve displays a softening behavior up to fracture, which occurs

at 850 ± 4 MPa stress and 14 ± 1% strain. The softening-like behavior is similar to that

observed for the Al-Ni-La alloy discussed in section 4.2. The stress-strain curve of the

sample HP at 723 K exhibits lower yield and compressive strengths (780 ± 10 MPa and

820 ± 2 MPa) as compared to the sample hot-pressed at 673 K. However, the fracture

strain is remarkably larger, reaching a value of 30 ± 2% at 690 ± 5 MPa stress.

Figure 4.3.5 Compression true stress-true strain curves of samples of as-atomized Al90.4Y4.4Ni4.3Co0.9 powder consolidated by hot pressing at 673 and 723 K.

The higher stress level of the sample hot-pressed at 673 K with respect to the

sample HP at 723 K is most likely due to the presence of the residual amorphous phase, as

shown by the corresponding XRD pattern in Figure 4.3.1 and the SEM micrograph in

Figure 4.3.3(b). Due to the incomplete crystallization of the amorphous phase, the sample

consolidated at 673 K displays a lower amount of crystallized fcc Al. As a result, the

fracture strain of the sample HP at 673 K is reduced compared to the sample HP at 723 K

that, in contrast, is fully crystallized and is not characterized by the hindrance of the

dislocations movement induced by the residual amorphous phase.

The current results clearly demonstrate that Al-based materials characterized by

high strength combined with considerable plastic strain can be produced through the

combined devitrification and consolidation of glassy precursors. The room temperature

mechanical properties of the materials investigated in the previous sections have been

tested in compression because the limited dimensions of the HP and SPS samples do not

94

permit accurate tensile tests. However, for the full evaluation of the mechanical behavior

of a material, tensile stress-strain data are necessary. Accordingly, preliminary tensile

results have been obtained for Al90.4Y4.4Ni4.3Co0.9 samples consolidated by hot pressing

followed by hot extrusion at 723 K. Through the extrusion of the HP samples, specimens

with dimensions suitable for standard tensile tests can be produced.

Figure 4.3.6 SEM micrographs of the Al90.4Y4.4Ni4.3Co0.9 samples consolidated by hot pressing followed by hot extrusion at 723 K.

Similarly to the sample HP at 723 K, the XRD pattern of the sample hot-extruded at

723 K (Figure 4.3.1) reveals diffraction peaks from Al, Al4Ni3, Al3Y and Al9Co2 phases

without any visible residual amorphous phase. The relative amounts of the phases is

slightly different for the sample hot-extruded at 723 K, which shows a larger amount of

Al3Y and a lower amount of Al4Ni3 compared to the sample that was hot-pressed at 723 K

(Figure 4.3.1). In addition, the diffraction peaks belonging to fcc Al are sharper, indicative

of a larger grain size. The grain size is about 150 nm for extruded sample as compared to

≈ 120 nm for the HP samples according to Scherer formula. This is due to the longer

dwelling time at 723 K of the double-step hot-pressing/hot-extrusion process (20 minutes)

compared to the single hot-pressing process (3 minutes). Figure 4.3.6 shows SEM

micrographs of the sample hot-extruded at 723 K, which reveals a microstructure

consisting of bright rod-like particles (the intermetallic compounds) distributed in dark-

gray background (fcc Al). In contrast to the material HP at 723 K (Figure 4.3.4), the

particles do not maintain the spherical shape. Also, no particle boundaries can be observed.

Due to the extrusion step, the nanocrystalline intermetallic particles are well-distributed

within the fcc Al giving rise to a homogeneous composite microstructure. In hot pressing

the powder particles shows spherical shape consists of nanocrystalline particles in Al

matrix, bright interfaces and congregation of Al at the triangular areas. This type of mixed

95

structure is completely sheared due to extrusion and results two phase structure of

nanocrystalline particles distributed in Al matrix without any boundaries.

Figure 4.3.7 shows the compressive and tensile stress-strain curves of the sample

hot extruded at 723 K. The hot-extruded Al90.4Y4.4Ni4.3Co0.9 material exhibits compressive

yield and fracture stresses of about 490 ± 5 MPa and 645 ± 3 MPa combined with 15 ± 1%

plastic strain. These values are smaller than those observed for the sample HP at 723 K

(Figure 4.3.5), probably because of the different microstructure induced by the additional

extrusion step. When tested in tension, the extruded specimen exhibits yield and fracture

stresses of about 480 ± 5 MPa and 565 ± 2 MPa, which are rather similar to the

compressive results. On the other hand, the tensile ductility compared to compression is

remarkably reduced (4 ± 1%), presumably because of the residual porosity (~ 1%) which is

more critical in the tensile mode for reducing the plastic strain through crack formation and

propagation than in the compressive mode [German 2008].

Figure 4.3.7 Room temperature tensile and compressive true stress-true strain curves for the samples consolidated by hot extrusion carried at 723 K of as-atomized Al90.4Y4.4Ni4.3Co0.9 powder.

The small difference between compressive and tensile strength observed in this

work clearly indicates that compression tests can be successfully used for a preliminary

evaluation of the strength of the consolidated materials. In addition, these preliminary

tensile results further demonstrate the validity of the combined devitrification and

consolidation of glassy precursors as a suitable method for the production of Al-based

materials characterized by high strength and considerable plastic deformation.

96

Chapter 5: Conclusions and outlook

Amorphous, partially amorphous and nanocrystalline Al-based alloys have been

attracting widespread attention as potential candidates for structural as well as functional

applications due to their high strength combined with low density. Although these

materials exhibit improved mechanical properties compared to conventional Al-based

crystalline alloys, the maximum scale of the products is limited to a thickness of less than

100 micrometers due to their relatively low glass-forming ability. In general, Al-based

metallic glasses and nanostructured materials with high Al content can only be obtained by

melt-spinning in the shape of ribbons or by gas atomization in the form of powder. This

limitation has prevented a wide extension of application fields of the Al-based amorphous

and nanocrystalline alloys even despite their excellent mechanical properties. To overcome

this limitation, powder metallurgical methods, such as gas atomization or mechanical

alloying followed by powder consolidation, can be employed to create bulk Al-based

samples with the desired microstructure. Accordingly, in this work novel bulk Al-based

alloys with high content of Al have been produced by powder metallurgy methods from

amorphous and partially amorphous materials. The present work focused on three specific

aspects:

(1) Production and characterization of Al-based amorphous and partially

amorphous powders with high Al content (> 80 at.%).

(2) Consolidation of the powder precursors into bulk samples with the

desired microstructure by different techniques.

(3) Microstructural characterization and mechanical property evaluation of

the consolidated bulk specimens.

Different processing routes, including mechanical alloying of elemental powder

mixtures, pulverization of melt-spun glassy ribbons and gas atomization, have been used

for the production of the Al-based powders. Although the mechanically-alloyed

Al85Y8Ni5Co2 powders reveal promising results in terms of glass formation and stability,

the milling time needed for amorphization is extremely long and the production yields poor

output, which drastically limits the use of mechanical alloying for the production of Al-

based powders with high Al contents. A better approach in terms of output of the powders

is the production of glassy powders by controlled pulverization of melt-spun ribbons. In

order to retain their glassy structure and to avoid sticking of the material to the milling

97

tools due to the high ductility of the ribbons, proper milling conditions have to be used (i.e.

interval-milling at a low intensity, corresponding to a rather low kinetic energy, performed

at the cryogenic temperature). Although this technique allows the production of larger

quantities of high quality material (i.e. constant composition and microstructure combined

with low contamination levels) compared to mechanical alloying, the procedure is rather

complex and requires the detailed characterization of every single melt-spun ribbon. On

the other hand, gas atomization offers the possibility to easily produce large quantities of

powders with homogeneous properties (e.g. structure and thermal stability) along with a

uniform size distribution of particles. Therefore, gas atomization is the best choice for the

production of Al-based amorphous and partially amorphous powders as precursors for the

subsequent consolidation step.

Materials in powder form have to be consolidated to achieve dense bulk specimens.

Consolidation of metastable phases, such as amorphous and nanocrystalline materials, is

not a trivial process and often results in undesirable microstructural transformations (e.g.

crystallization and grain coarsening), or insufficient particle bonding. These characteristics

severely limit the consolidation parameters that can be used and, as a result, temperature,

pressure and the time span of the consolidation process have to be adjusted carefully in

order to achieve a balance between good densification and desired microstructure. For that

reason, the crystallization behavior and the temperature dependence of the viscosity have

been studied in detail in order to optimize the processing conditions and to select the

proper consolidation parameters.

Following their characterization, the Al-based powders have been consolidated into

bulk specimens by hot pressing (HP), hot extrusion and spark plasma sintering (SPS) and

their microstructure and mechanical properties have been investigated. The results indicate

that the mechanical properties of the consolidated samples can be varied within a wide

range of strength and ductility depending on the microstructure and the consolidation

techniques used.

Single-phase amorphous bulk Al85Y8Ni5Co2 specimens were produced by hot

pressing of the pulverized ribbons. Room temperature compression tests of the single-

phase glass reveal low strength and no ductility due to the residual porosity of the

consolidated specimen. No extrusion of the single-phase glass was possible at temperatures

below the crystallization temperature due to the insufficient viscous flow of the SCL in the

present consolidation conditions. In order to reach a higher density and, consequently, to

improve the ductility of the samples, the milled amorphous ribbons were blended with

98

different volume fractions of elemental fcc Al to produce glass-matrix composites. The

resulting powders were then consolidated by hot pressing followed by hot extrusion. When

50 vol.% Al is added, the maximum stress is found to be 295 MPa, therefore, decreased

with respect to the single-phase amorphous specimen (400 MPa). However, the strain at

maximum stress is about 7 %. The material containing 70 vol.% Al exhibits a maximum

stress of 255 MPa and a strain at maximum stress is about 10 %. These results indicate that

glass-reinforced Al-based composites with high strength combined with considerable

ductility can be produced by powder metallurgy methods. The mechanical properties of the

glass-reinforced composites can be modeled by using the iso-stress Reuss model, which

allows the prediction of the mechanical properties of a composite from the volume-

weighted averages of the components properties.

Higher strength levels combined with good plastic deformation at room

temperature can be achieved by the combined devitrification and consolidation of gas-

atomized amorphous and partially amorphous precursors. For this purpose, Al-based gas-

atomized powders (GAP) with compositions Al84Gd6Ni7Co3, Al87Ni8La5 and

Al90.4Y4.3Ni4.4Co0.9 have been used as precursors for the production of high-strength bulk

samples.

The results on hot-pressing of the Al84Gd6Ni7Co3 powder indicate that

consolidation into highly-dense bulk samples cannot be achieved without extended

crystallization of the material. Nevertheless, crystallization during consolidation is not

detrimental and leads to bulk samples with a remarkably high strength of about 1500 MPa,

which is three times larger than the conventional high-strength Al-based alloys.

Investigation of the sintering behavior reveals that preferential diffusion of Al toward the

pore regions occurs during hot pressing, leading to the filling of the triangular areas

between the particles by ultra fine-grained Aluminum. This, together with the formation of

rod-like intermetallics, leads to bulk samples characterized by high strength combined with

a limited but distinct plastic deformability (3.5 - 4%).

To further test the effectiveness of powder consolidation as a method for the

production of Al-based materials with enhanced mechanical properties through the

combined crystallization and consolidation of glassy precursors, Al87Ni8La5 gas-atomized

powders have been consolidated by SPS above their crystallization temperature. Spark

plasma sintering leads to highly dense bulk specimens with a multi-phase structure

consisting of fcc-Al together with Al11La3 and Al3Ni intermetallic compounds. The

consolidated bulk material exhibits high compression strength of 930 MPa together with

99

plastic strain exceeding 25 %. The high deformation capability is most likely due to the

formation of a microstructure consisting of a network of ultra fine-grained Al reinforced

with nm-scale intermetallic particles. Within this structure, the fcc Al regions are not

confined and, as a result, the continuous network of fcc Al may allow the movement of

dislocations, explaining the remarkable plastic deformation with respect to the

Al84Gd6Ni7Co3 bulk sample.

In order to investigate the effect of high Al content on the mechanical behavior of

the consolidated samples, Al90.4Y4.4Ni4.3Co0.9 gas-atomized powder have been consolidated

by hot pressing. The bulk samples display remarkable mechanical properties, namely, high

compression strength ranging between 820 and 925 MPa combined with plastic strain in

the range 14 – 30%. Strength and plastic strain of the hot-pressed samples are strictly

linked with their microstructure. Higher strength and reduced plasticity are related to the

presence of a residual amorphous phase, which may hinder the dislocations movement

within the Al phase. On the other hand, reduced strength but enhanced plastic deformation

is a result of the complete crystallization of the glass and of the formation of additional fcc

Al from the residual amorphous phase. In addition, preliminary tensile tests for the

Al90.4Y4.3Ni4.4Co0.9 alloy consolidated by hot pressing followed by hot extrusion reveal

promising tensile properties (tensile strength 565 MPa and 4% ductility) comparable to

commercial high-strength Al-based alloys.

The results presented in this thesis clearly indicate that powder metallurgy, i.e.

powder synthesis and consolidation, is a particularly suitable method for the production of

Al-based materials characterized by high strength combined with considerable plastic

strain. The combined devitrification and consolidation of glassy precursors into high-

strength deformable bulk samples promise a new route for the development of novel and

innovative high-performance Al-based materials for eco-friendly transport applications.

However, some aspects have still to be investigated. For example, the mechanism

responsible for the preferential diffusion of Al toward the pore regions has to be fully

clarified and supplementary investigations, such as additional studies on the effect of the

glass crystallization and related atomic diffusion on the sintering behavior and resulting

microstructure, are required to clarify this aspect. Also, systematic investigations on the

tensile properties of these novel Al-based alloys have to be carried out. In addition, fatigue,

wear and corrosion properties of the consolidated materials, which are crucial aspects for

any potential commercial application, have to be fully evaluated.

100

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Selbständigkeitserklärung

Hiermit erkläre ich, dass ich die vorliegende Arbeit ohne die unzulässige Hilfe Dritter und

ohne Benutzung anderer als der angegebenen Hilfsmittel angefertigt habe, die aus fremden

Quellen direkt oder indirekt übernommenen Gedanken sind als solche kenntlich gemacht. Die

Arbeit wurde bisher weder im Inland noch im Ausland in gleicher oder ähnlicher Form einer

anderen Prüfungsbehörde vorgelegt.

Die Arbeit „Production of high strength Al-based alloys by consolidation of amorphous and

partially amorphous powders“, vorgelegt von Kumar ‚Babu Surreddi, wurde unter Betreuung

von Prof. Dr. Jürgen Eckert am Institut für Komplexe Materialien (IKM) des Leibniz Institut

für Festkörper- und Werkstoffforschung Dresden e.V. (IFW Dresden) angefertigt.

Ich erkenne hiermit die Promotionsordnung der Technischen Universität Dresden an.

Dresden, den 03 June 2011

(Kumar Babu Surreddi)

List of Publications

[1-23] [1] S. Scudino, K. B. Surreddi and J. Eckert, Mechanical properties of cold-rolled

Zr60Ti5Ag5Cu12.5Ni10Al7.5 metallic glass, Physica Status Solidi A-Applications and

Materials Science, 207 (2010) 1118-1121.

[2] S. Scudino, K. B. Surreddi, G. Wang and J. Eckert, Enhanced plastic deformation of

Zr41.2Ti13.8Cu12.5Ni10Be22.5 bulk metallic glass by the optimization of frictional

boundary restraints, Scripta Materialia, 62 (2010) 750-753.

[3] V.C. Srivastava, K. B. Surreddi, S. Scudino, M. Schowalter, V. Uhlenwinkel, A.

Schulz, J. Eckert, A. Rosenauer and H.W. Zoch, Microstructure and mechanical

properties of partially amorphous Al85Y8Ni5Co2 plate produced by spray forming,

Materials Science and Engineering A, 527 (2010) 2747-2758.

[4] K. B. Surreddi, S. Scudino, M. Sakaliyska, K.G. Prashanth, D.J. Sordelet and J.

Eckert, Crystallization behavior and consolidation of gas-atomized Al84Gd6Ni7Co3

glassy powder, Journal of Alloys and Compounds, 491 (2010) 137-142.

[5] S. Scudino, M. Sakaliyska, K. B. Surreddi, F. Ali and J. Eckert, Structure and

mechanical properties of Al-Mg alloys produced by copper mold casting, Journal of

Alloys and Compounds, 504 (2010) S483-S486.

[6] S. Scudino, K. B. Surreddi, M.S. Khoshkhoo, M. Sakaliyska, G. Wang and J. Eckert,

Improved Room Temperature Plasticity of Zr41.2Ti13.8Cu12.5Ni10Be22.5 Bulk Metallic

Glass by Channel-Die Compression, Advanced Engineering Materials, 12 (2010) 8

(online).

[7] S. Scudino, F. Ali, K. B. Surreddi, K.G. Prashanth, M. Sakaliyska and J. Eckert, Al-

based metal matrix composites reinforced with nanocrystalline Al-Ti-Ni particles,

Journal of Physics: Conference Series, 240 (2010) 012154.

[8] K. B. Surreddi, V.C. Srivastava, S. Scudino, M. Sakaliyska, V. Uhlenwinkel, J.S.

Kim and J. Eckert, Production of high-strength Al85Y8Ni5Co2 bulk alloy by spark

plasma sintering, Journal of Physics: Conference Series, 240 (2010) 012155.

[9] V.C. Srivastava, K. B. Surreddi, S. Scudino, M. Schowalter, V. Uhlenwinkel, A.

Schulz, A. Rosenauer, H.W. Zoch and J. Eckert, Spray forming of bulk Al85Y8Ni5Co2

with co-existing amorphous, nano- and micro-crystalline structures, Transactions of

the Indian Institute of Metals, 62 (2009) 331-335.

[10] M. Sakaliyska, S. Scudino, H.V. Nguyen, K. B. Surreddi, B. Bartusch, F. Ali, J.S.

Kim and J. Eckert, Consolidation and mechanical properties of mechanically alloyed

Al-Mg powders, Advanced Intermetallic-Based Alloys for Extreme Environment and

Energy Applications, 1128 (2009) 299-304.

[11] S. Scudino, M. Sakaliyska, K. B. Surreddi and J. Eckert, Solid-state processing of Al-

Mg alloys Journal of Physics: Conference Series, 144 (2009) 12019-12019.

[12] K. B. Surreddi, S. Scudino, H.V. Nguyen, K. Nikolowski, M. Stoica, M. Sakaliyska,

J.S. Kim, T. Gemming, J. Vierke, M. Wollgarten and J. Eckert, Spark plasma sintering

of gas atomized Al87Ni8La5 amorphous powder Journal of Physics: Conference Series,

144 (2009) 12079-12079.

[13] K. Nikolowski, S. Scudino, M. Stoica, K. B. Surreddi, J. Das and J. Eckert, Stress-

induced martensitic transformation in a Ti45Zr38Al17 cast rod Journal of Physics:

Conference Series, 144 (2009) 12090-12090.

[14] S. Scudino, M. Sakaliyska, K. B. Surreddi and J. Eckert, Mechanical alloying and

milling of Al-Mg alloys, Journal of Alloys and Compounds, 483 (2009) 2-7.

[15] S. Scudino, S. Venkataraman, M. Stoica, K. B. Surreddi, S. Pauly, J. Das and J.

Eckert, Consolidation and mechanical properties of ball milled Zr50Cu50 glassy ribbons,

Journal of Alloys and Compounds, 483 (2009) 227-230.

[16] S. Scudino, K. B. Surreddi, H.V. Nguyen, G. Liu, T. Gemming, M. Sakaliyska, J.S.

Kim, J. Vierke, M. Wollgarten and J. Eckert, High-strength Al87Ni8La5 bulk alloy

produced by spark plasma sintering of gas atomized powders, Journal of Materials

Research, 24 (2009) 2909-2916.

[17] S. Scudino, G. Liu, M. Sakaliyska, K. B. Surreddi and J. Eckert, Powder metallurgy

of Al-based metal matrix composites reinforced with beta-Al3Mg2 intermetallic

particles: Analysis and modeling of mechanical properties, Acta Materialia, 57 (2009)

4529-4538.

[18] K.G. Prashanth, S. Scudino, K. B. Surreddi, M. Sakaliyska, B.S. Murty and J. Eckert,

Crystallization kinetics of Zr65Ag5Cu12.5Ni10Al7.5 glassy powders produced by ball

milling of pre-alloyed ingots, Materials Science and Engineering A, 513-14 (2009)

279-285.

[19] S. Scudino, P. Donnadieu, K. B. Surreddi, K. Nikolowski, M. Stoica and J. Eckert,

Microstructure and mechanical properties of Laves phase-reinforced Fe-Zr-Cr alloys,

Intermetallics, 17 (2009) 532-539.

[20] S. Scudino, G. Liu, K.G. Prashanth, B. Bartusch, K. B. Surreddi, B.S. Murty and J.

Eckert, Mechanical properties of Al-based metal matrix composites reinforced with

Zr-based glassy particles produced by powder metallurgy, Acta Materialia, 57 (2009)

2029-2039.

[21] V.C. Srivastava, K. B. Surreddi, V. Uhlenwinkel, A. Schulz, J. Eckert and H.W.

Zoch, Formation of Nanocrystalline Matrix Composite during Spray Forming of

Al83La5Y5Ni5Co2, Metallurgical and Materials Transactions A, 40A (2009) 450-461.

[22] S. Scudino, M. Sakaliyska, M. Stoica, K. B. Surreddi, F. Ali, G. Vaughan, A.R.

Yavari and J. Eckert, In-situ X-ray diffraction of mechanically milled beta-Al3Mg2

powders, Physica Status Solidi-Rapid Research Letters, 2 (2008) 272-274.

[23] S. Scudino, K. B. Surreddi, S. Sager, M. Sakaliyska, J.S. Kim, W. Loser and J.

Eckert, Production and mechanical properties of metallic glass-reinforced Al-based

metal matrix composites, Journal of Materials Science, 43 (2008) 4518-4526.