the pennsylvania state university the graduate school

The Pennsylvania State University

The Graduate School

College of Earth and Mineral Sciences

METAL ORGANIC CHEMICAL VAPOR DEPOSITION OF ENVIRONMENTAL

BARRIER COATINGS FOR THE INHIBITION OF SOLID DEPOSIT FORMATION

FROM HEATED JET FUEL

A Dissertation in

Energy and Geo-Environmental Engineering

by

Arun Ram Mohan

©2011 Arun Ram Mohan

Submitted in Partial Fulfillment

of the Requirements

for the Degree of

Doctor of Philosophy

May 2011

ii

The dissertation of Arun Ram Mohan was reviewed and approved* by the following:

Semih Eser

Professor of Energy and Geo-Environmental Engineering

Dissertation Advisor

Chair of Committee

Ljubisa R. Radovic

Professor of Energy and Mineral Engineering

André L. Boehman

Professor of Fuel Science and Materials Science and Engineering

Coray M. Colina

Corning Faculty Fellow

Associate Professor of Materials Science and Engineering

Yaw D. Yeboah

Professor and Department Head of Energy and Mineral Engineering

*Signatures are in file in the Graduate School

iii

ABSTRACT

Solid deposit formation from jet fuel compromises the fuel handling system of an aviation

turbine engine and increases the maintenance downtime of an aircraft. The deposit formation

process depends upon the composition of the fuel, the nature of metal surfaces that come in

contact with the heated fuel and the operating conditions of the engine. The objective of the

study is to investigate the effect of substrate surfaces on the amount and nature of solid deposits

in the intermediate regime where both autoxidation and pyrolysis play an important role in

deposit formation. A particular focus has been directed to examining the effectiveness of barrier

coatings produced by metal organic chemical vapor deposition (MOCVD) on metal surfaces for

inhibiting the solid deposit formation from jet fuel degradation.

In the first part of the experimental study, a commercial Jet-A sample was stressed in a

flow reactor on seven different metal surfaces: AISI316, AISI 321, AISI 304, AISI 347, Inconel

600, Inconel 718, Inconel 750X and FecrAlloy. Examination of deposits by thermal and

microscopic analysis shows that the solid deposit formation is influenced by the interaction of

organosulfur compounds and autoxidation products with the metal surfaces. The nature of metal

sulfides was predicted by Fe-Ni-S ternary phase diagram. Thermal stressing on uncoated

surfaces produced coke deposits with varying degree of structural order. They are hydrogen-rich

and structurally disordered deposits, spherulitic deposits, small carbon particles with relatively

ordered structures and large platelets of ordered carbon structures formed by metal catalysis.

In the second part of the study, environmental barrier coatings were deposited on tube

surfaces to inhibit solid deposit formation from the heated fuel. A new CVD system was

configured by the proper choice of components for mass flow, pressure and temperature control

in the reactor. A bubbler was designed to deliver the precursor into the reactor for the deposition

iv

of metal and metal oxide functional coatings by MOCVD. Alumina was chosen as a candidate

for metal oxide coating because of its thermal and phase stability. Platinum was chosen as a

candidate to utilize the oxygen spillover process to maintain a self-cleaning surface by oxidizing

the deposits formed during thermal stressing. Two metal organic precursors, aluminum

trisecondary butoxide and aluminum acetylacetonate, were used as precursors to coat tubes of

varying diameters. The morphology and uniformity of the coatings were characterized by

electron microscopy and energy-dispersive x-ray spectroscopy. The coating was characterized by

x-ray photoelectron spectroscopy to obtain the surface chemical composition. This is the first

study conducted to examine the application of MOCVD to coat internal surfaces of tubes with

varying diameters.

In the third part of the study, the metal oxide coatings, alumina from aluminum

acetylacetonate, alumina from aluminum trisecondary butoxide, zirconia from zirconium

acetylacetonate, tantalum oxide from tantalum pentaethoxide and the metal coating, platinum

from platinum acetylacetonate were deposited by MOCVD on AISI304. The chemical

composition and the surface acidity of the coatings were characterized by x-ray photoelectron

spectroscopy. The morphology of the coatings was characterized by electron microscopy. The

coated substrates were tested in the presence of heated Jet-A in a flow reactor to evaluate their

effectiveness in inhibiting the solid deposit formation. All coatings inhibited the formation of

metal sulfides and the carbonaceous solid deposits formed by metal catalysis. The coatings also

delayed the accumulation of solid carbonaceous deposits. In particular, it has been confirmed

that the surface acidity of the metal oxide coatings affects the formation of carbonaceous

deposits. Bimolecular addition reactions promoted by the Brønsted acid sites appear to lead to

the formation of carbonaceous solid deposits depending on the surface acidity of the coatings.

v

In the last part of the study, the residual carbon was incorporated in the zirconia coating by

deposition with and without oxygen. As carbon surface is less active towards coke deposition,

presence of residual carbon in the coating was expected to reduce its activity towards carbon

deposition. The residual carbon in the coating was characterized by Raman spectroscopy and

thermal analysis. However, it has been observed that residual carbon in the coating beyond a

certain concentration compromises the integrity of the coating during the process of cooling the

substrate from deposition temperature to room temperature. It has been found that residual

carbon in the zirconia coating does not appear to affect the activity of the surface towards carbon

deposition.

vi

Table of Contents

List of Tables ix

List of Figures x

Acknowledgements xv

Chapter 1. Introduction 1

1.1 Background 1

1.1.1 Liquid phase autoxidation of hydrocarbons 1

1.1.2 Deposit formation during autoxidation of hydrocarbons 4

1.1.2.1 Effects of hydrocarbon structure 4

1.1.2.2 Effects of dissolved oxygen 7

1.1.2.3 Effects of sulfur compounds 8

1.1.2.4 Effects of nitrogen- and oxygen-containing compounds 9

1.1.2.5 Effects of antioxidants 9

1.1.2.6 Synergism between natural and synthetic antioxidants 10

1.1.2.7 Effects of natural antioxidants in autoxidation of neat and blended fuels 12

1.1.2.8 Effects of surface catalysis on the liquid phase autoxidation of hydrocarbons 13

1.1.2.8a Effects of metals 13

1.1.2.8b Effects of metal oxides 15

1.1.2.9 Autoxidation of jet fuels and deposit formation 16

1.1.3 Factors affecting deposits under pyrolytic conditions 20

1.1.4 Solid deposit formation from jet fuel under pyrolytic conditions 21

1.1.4.1 Carbonaceous mesophase 21

1.1.4.2 Filamentous carbon 23

1.1.4.3 Spherulitic deposits 24

1.1.4.4 Pyrolytic carbon 25

1.1.4.5 Metal sulfides 25

1.2 Objectives of the thesis 27

1.3 Organization of the thesis 28

1.4 References 29

Chapter 2. Analysis of Carbonaceous Solid Deposits from Thermal Oxidative Stressing of Jet-A

Fuel on Iron and Nickel-based Alloy Surfaces 34

2.1 Abstract 34

2.2 Introduction 34

2.3 Experimental Section 36

2.3.1 Thermal stressing experiments 36

2.3.2 Characterization of carbon deposits 37

2.4 Results and Discussion 38

2.4.1 Amount of solid carbon deposited on different metal substrates 39

2.4.2 TPO and FESEM analysis of deposits on various substrates 40

2.5 Conclusions 48

2.6 References 50

vii

Chapter 3. Environmental Barrier Coatings by MOCVD on Tube Surfaces to Inhibit Carbon

Deposition 61

3.1 Background 61

3.2 Coating process for EBCs 63

3.2.1 Plasma spray deposition 63

3.2.2 Electron beam physical vapor deposition 66

3.2.3 Electrodeposition process 68

3.2.4 Chemical vapor deposition 70

3.2.5 Effect of process variables on MOCVD and properties of coatings 71

3.2.6 Influence of process parameters in the stress induced in coatings 73

3.2.7 Coating precursors 74

3.2.8 Configuration of the MOCVD experimental set-up for coating tubes 75

3.3 Experimental Procedure 76


3.4.1 Characterization of alumina coatings from aluminum trisecondary butoxide 77

3.4.2 Characterization of alumina coatings from aluminum acetylacetonate 84

3.4.3 Temperature-Programmed Oxidation of residual carbon in alumina coatings 88

3.4.4 Characterization of platinum deposited on alumina coating 89

3.5 Conclusions 91

3.6 References 93

Chapter 4. Effectiveness of Low-Pressure MOCVD Coatings on Metal Surfaces for the

Mitigation of Fouling from Heated Jet Fuel 112

4.1 Abstract 112

4.2 Introduction 113

4.3 Experimental section 114

4.3.1 MOCVD experimental set-up for foil coatings 114


4.3.3 Characterization of coatings and carbon deposits 117


4.4.1 Morphology and spectroscopic characterization of coated substrates 120

4.4.2 Analysis of Jet-A sample and TPO of the deposits on coated & uncoated

substrates 125

4.5 Conclusions 132

4.6 References 134

Chapter 5. Characterization of Zirconia Coatings Deposited by MOCVD and Their

Effectiveness in Inhibiting Solid Deposition from Jet Fuel 150

5.1 Abstract 150

5.2 Introduction 150

5.3 Experimental section 151

5.3.1 MOCVD experimental set-up for zirconia coating 151


5.3.3 Characterization of coatings and carbon deposits 153

viii


5.4.1 Morphology of zirconia coating 155

5.4.2 Raman spectra of zirconia coatings 156

5.4.3 Infrared spectrum of zirconia coatings 157

5.4.4 X-ray photoelectron spectroscopy of residual carbon in the zirconia coating 159

5.4.5 Temperature–Programmed Oxidation 160

5.5 Conclusions 162

5.6 References 163

Chapter 6. Conclusions, Summary, and Recommendations for Future Work 170

6.1 Conclusions 170

6.2 Summary 171

6.3 Recommendations for Future Work 174

Appendix

Appendix A. Repeatability Data for TPO Profile of Solid Deposits on Substrates 175

Appendix B. Calculations for the Amount of Carbon Deposits on Various Substrate 182

Surfaces

ix

List of Tables

Table 2.1. Elemental composition of alloys. (Goodfellow Ltd) 53

Table 2.2. Calculated atomic percentage of Fe, Ni, and S on the alloys after 5 hours of thermal

stressing. 53

Table 3.1. Benefits and limitations of various deposition methods 97

Table 3.2. Description of the components used in the MOCVD system 98

Table 3.3. Growth conditions of the coating and their respective characteristics 98

Table 3.4. Composition of the coating by X-ray photoelectron spectroscopy 99

Table 3.5. Concentration of various functional groups under C 1s peak 99

Table 4.1. Conditions used to deposit metal oxide coatings by MOCVD 138

Table 4.2. Relative areas (%) of the deconvoluted peaks under the N1s scan 138

x

List of Figures

Figure 1.1. Effect of hydrocarbon structure in deposit formation 7

Figure 1.2. Chemical and physical process in thermal deposition 18

Figure 1.3. Schematic showing the growth mechanism of Fe-S nanostructure 26

Figure 2.1. Flow reactor setup for thermal stressing experiment 54

Figure 2.2. (a) GC-MS chromatogram of Jet-A showing the composition of the fuel (b) PFPD

chromatogram showing sulfur compounds. 54-55

Figure 2.3. Carbon deposits on different metal surfaces from Jet-A at 350 °C, 500 psig with a

fuel flow rate of 1 mL/min for 5 h 55

Figure 2.4. Fe-Ni-S ternary phase diagram at 400 ºC. 56

Figure 2.5. (a) FESEM image and (b) TPO profile of the deposits formed on AISI 316 from Jet-

A at 350 °C and 500 psig for 5 h. 56



Figure 2.7. (a) FESEM image (b) X-ray diffractogram and (c) TPO profile of the deposits

formed on AISI 304 from Jet-A at 350 °C and 500 psig for 5 h. 57



Figure 2.9. (a) FESEM image and (b) TPO profile of the deposits formed on FecrAlloy from Jet-


Figure 2.10. (a) FESEM image and (b) TPO profile of the deposits formed on Inconel 600 from

Jet-A at 350 °C and 500 psig for 5 h. 59


Jet-A at 350 °C and 500 psig for 5 h. 59

Figure 2.12. (a) FESEM image and (b) TPO profile of the deposits formed on Inconel 750X

from Jet-A at 350 °C and 500 psig for 5 h. 60

Figure 2.13. (a) FESEM image (b) TPO profile of the deposits formed on Silicon from Jet-A at

350 °C and 500 psig for 5 h. 60

Figure 3.1. Atmospheric plasma spray method for TBCs. 64

xi

Figure 3.2. Schematic representation of the MOCVD setup used for coating tubes 100

Figure 3.3. Structure of the metalorganic precursors. 100

Figure 3.4. (a) Photograph of the aluminum block used for heating the fuel injector (b)

Photographs of a fuel injector nozzle (c) Photographs of a fuel injector before and after the

attachment of the fitting at the tip (d) Photograph showing the injector tip before (right) and after

deposition (left). 101

Figure 3.5. (a) Photograph of an uncoated ¼‖electropolished stainless steel tube heated to 500

°C in the presence of argon (b) Photograph of the coating ATSB-A at 400 °C sectioned into two

halves (c) Photograph of the coating ATSB-C deposited at 450 °C. 102

Figure 3.6. (a) SEM micrograph of alumina coating ATSB-A deposited at 400 °C. EDX

elemental map of (b) aluminum and (c) oxygen (d) carbon (e) EDX signal corresponding to an

elemental composition of 39% of aluminum and 61% of oxygen. 103

Figure 3.7. (a) Cross-sectional SEM micrograph of alumina coating ATSB-A deposited from

aluminum trisecondary butoxide at TS = 400 °C and TB = 160 °C. EDX elemental map of (b)

aluminum (c) oxygen (d) Iron. 104

Figure 3.8. (a) SEM micrograph of alumina coating ATSB-B deposited from aluminum

trisecondary butoxide at TR = 400 °C and TB = 132 °C. Elemental map of (b) Aluminum (c)

oxygen (d) carbon on the coating (e) EDX spectrum of the coating. 105

Figure 3.9 SEM micrographs of alumina coating ATSB-B deposited from aluminum

trisecondary butoxide at TS = 400 °C and TB = 132 °C. 106

Figure 3.10. SEM micrograph of alumina coating from aluminum trisecondary butoxide

evaporated at 132 °C and deposited at (a) 450 °C for ATSB-C (b) 500 °C for ATSB-D. 106

Figure 3.11. (a) Photograph of the coated tube AlacacA 106

Figure 3.12. (a) SEM micrograph of alumina coating AlacacA from aluminum acetylacetonate

sublimed at 138 °C and deposited at 400 °C (b) Elemental map of Aluminum (c) oxygen (d)

carbon (e) EDX spectrum of the coating. 107

Figure 3.13. (a) Photograph of the uncoated tube. (b) Photograph of coated tube AlacacB 108

Figure 3.14. (a) SEM micrograph of alumina coating AlacacB from aluminum acetylacetonate

sublimed at 138 °C and deposited at 500 °C. (b) Elemental map of Aluminum (c) oxygen (d)

carbon. 108

Figure 3.15. (a) XPS survey scan of thermally stressed alumina film AlacacB on AISI 304 after

temperature programmed oxidation from 100 °C – 900 °C and holding the coating at 900 °C for

5 minutes in UHP oxygen.(b) SEM micrograph of the coating at 30μm 109

xii

Figure 3.16. (a) Photograph of coated 1/8‖ tube (b) SEM micrograph of alumina coating

AlacacC from aluminum acetylacetonate sublimed at 138 °C and deposited at 500 °C after 4

hours at 30x. (c) Elemental map of Aluminum (d) oxygen and (f) carbon and XPS survey scan of

Alumina coating on the one eighth of an inch tube. 109-110

Figure 3.17. (a) SEM micrograph of alumina coating AlacacC from aluminum acetylacetonate

sublimed at 138 °C and deposited at 500 °C after 24 hours. (b) EDX spectrum of alumina coating

on the one eighth of an inch tube showing the presence of aluminum and oxygen. 110

Figure 3.18. (a) TPO of residual carbon in the alumina coating formed by the decomposition of

aluminum acetylacetonate. (b) High resolution scan of C 1s by XPS on the alumina coating

deposited at 450 °C. 111

Figure 3.19. Photograph of coated tube Pt-AlacacB. 111

Figure 3.20. (a) SEM micrograph of the coating Pt-AlacacB (b) High resolution scan for

platinum 4f from X-ray photoelectron spectroscopy. 111

Figure 4.1. Schematic diagram of the MOCVD set-up used for the deposition of metal oxide

coatings. 139

Figure 4.2. Flow reactor set-up for thermal stressing experiments with Jet-fuel. 139

Figure 4.3. Specimen preparation by FIB for TEM examination: SEM micrograph of the coating

after platinum deposition (a), the specimen cross-section after milling to form a wedge (b), and

the specimen fastened to the TEM grid (c). 140

Figure 4.4. An SEM image (a) and an AFM images (b) of the blank substrate AISI304. 140

Figure 4.5. An SEM image (a), an AFM image (b), a cross-sectional TEM image (c), a

diffraction pattern (d), and a high resolution scan for N1s after pyridine adsorption on alumina

coating from aluminum trisecondary butoxide (ATSB) (e). 141

Figure 4.6. An SEM images (a), an AFM image (b), a cross-sectional TEM image (c), a

diffraction pattern (d), a high resolution scan for N1s after pyridine adsorption on alumina

coating from aluminum acetylacetonate (Alacac) (e). 142

Figure 4.7. An SEM image (a), an AFM image (b), a cross-sectional TEM image (c), and a

diffraction pattern (d) of the Zirconia coating from zirconium acetylacetonate. 143

Figure 4.8. An SEM image (a) and an AFM image (b) of the Tantalum oxide coating from

tantalum pentaethoxide. 143

Figure 4.9. A GC-MS chromatogram of Jet-A with marked peaks for n-alkanes (a), and PFPD

chromatogram with identified sulfur compounds (b). 144

xiii

Figure 4.10. Amount of carbon deposits on metal and coated surfaces from Jet-A at 350 °C, 500

psig with a fuel flow rate of 1 mL/min for 5 h. 145

Figure 4.11. An SEM image (a), a X-ray diffractogram (b), and a TPO profile of deposits (c)

formed on AISI304 from Jet-A at 350 °C and 500 psig for 5 h. 146

Figure 4.12. An SEM image (a) and a TPO profile of carbonaceous deposits (b) from Jet-A at

350 °C and 500 psig for 5 h (b) on alumina coating from aluminum trisecondary butoxide

(ATSB). 146


350 °C and 500 psig for 5 h on alumina coating from aluminum acetylacetonate (Alacac). 147


350 °C and 500 psig for 5 h on Zirconia coating. 147


350 °C and 500 psig for 5 h on tantalum oxide coating. 148

Figure 4.16. An SEM image (a), a TPO profile of carbonaceous deposits from Jet-A at 350 °C

and 500 psig for 5 h on platinum coating (b), an EDX spectrum of the sample (c), an EDX map

of platinum (d), an EDX map of carbon (e) and an EDX map of sulfur (f). 148-149

Figure 5.1. Schematic of the MOCVD setup used for the deposition of zirconia coatings 165

Figure 5.2. (a) SEM of the coating deposited at 500 °C. EDX map of (b) zirconium (c) oxygen

and (d) carbon 165

Figure 5.3. Raman spectra of zirconia coating on AISI 304 (a) before TPO (b) after TPO 166

Figure 5.4. (a) Structure of the precursor zirconium acetylacetonate (b) Attenuated total

reflection infrared spectrum of zirconia coating on AISI 304 167

Figure 5.5. High resolution scan for C 1s in the zirconia coating deposited at 400 °C. 168

Figure 5.6. SEM micrograph of zirconia coating deposited at 450 °C. 168

Figure 5.7. TPO of residual carbon in the zirconia coatings deposited at 500 °C 169

Figure 5.8. (a) SEM micrograph of the zirconia coating deposited at 500 °C (b) TPO of the

zirconia coating after thermal stressing with Jet-A at 350 °C, 500 psig and 1 mL/min for 5 hours

carbon representing both carbonaceous deposits from Jet-A and carbon incorporated in the

coating from the precursor 169

Figure A.1. Reproducibility data for TPO: TPO profile of solid carbonaceous deposits on AISI

304 from Jet-A at 350 °C and 500 psig for 5 h. 175

xiv







Figure A.5. Reproducibility data for TPO: TPO profile of solid carbonaceous deposits on

FeCrAlloy from Jet-A at 350 °C and 500 psig for 5 h. 177


Inconel 600 from Jet-A at 350 °C and 500 psig for 5 h. 177


Inconel 718 from Jet-A at 350 °C and 500 psig for 5 h. 178


Inconel 750-X from Jet-A at 350 °C and 500 psig for 5 h. 178

Figure A.9. Reproducibility data for TPO: TPO profile of solid carbonaceous deposits from Jet-

A at 350 °C and 500 psig for 5 h on the alumina coating deposited from aluminum trisecondary

butoxide. 179

Figure A.10. Reproducibility data for TPO: TPO profile of solid carbonaceous deposits from

Jet-A at 350 °C and 500 psig for 5 h on the platinum coating deposited from platinum

acetylacetonate on AISI 304. 179


Jet-A at 350 °C and 500 psig for 5 h on the tantalum oxide coating deposited from tantalum

pentaethoxide on AISI 304. 180


Jet-A at 350 °C and 500 psig for 5 h on the alumina coating deposited from aluminum

acetylacetonate. 180


Jet-A at 350 °C and 500 psig for 5 h on the zirconia coating deposited from zirconium

acetylacetonate. 181

xv

Acknowledgements

I would like to thank my dear advisor Dr. Semih Eser, for his support and guidance

throughout my Ph.D thesis and making me a confident and an independent researcher over the

course of time. I would like to express my sincere gratitude to my committee members Dr.

Ljubisa Radovic, Dr. Andre Boehman and Dr. Coray Colina for serving in my committee and for

their valuable feedback. I appreciate Dr. Angela Lueking for her time and feedback during my

candidacy exam. I would like to thank Dr. Sarma Pisupati and Dr. Yaw Yeboah for providing me

financial support in the form of teaching assistantship during my graduate studies. I thank Rolls

Royce corporation, British Petroleum, Restek, Combustion Science and Engineering and the

EMS Energy Institute for their financial support at various times to carry out my research

activities for my Ph.D Thesis.

I would like to thank Dr Orhan Altin for leaving behind a good working thermal stressing

reactor that helped me to start my thesis work. I would like to thank Mr Ronald Wincek and

Glenn Decker for their valuable time in helping me to layout a completely new CVD

experimental system. Special thanks to Mr. Ronald Wincek for troubleshooting the TPO

instrument. I also would like to thank Mr. Ken Biddle, Mr. Bill Diehl and Mr. Bill Genet in the

EMS machine shop for their valuable engineering services that helped me to complete my

research activities for my sponsors on time. I thank Mr. Bob Hengstebeck of Material Research

Institute for teaching me XPS data analysis. I thank Dr. Tad Daniel and Mr. Vincent Bojan for

data acquisition from the XPS instrument. I thank Mr. John Cantolina for training me with the

FESEM. I thank Dr. Trevor Clark for his valuable assistance in FIB and TEM on my samples. I

thank Dr. Trevor Clark and Dr Nandakumar for sharing their career experience in various

organizations that helped to shape my values and professional objectives. I thank Dr. Joe Stit for

xvi

helping me in the data acquisition of Raman spectrum. I thank Dr. Josh Stapleton for the data

acquisition of IR spectrum. I thank Dr. Cigdam Shalaby in Dr. Song’s group for her assistance in

sulfur analysis.

I am indebted to my parents Dr. Ram Mohan and Mrs. Lakshmi who were my first

teachers for their extraordinary support in all aspects, values, guidance and prayers throughout

my life. I dedicate my Ph.D thesis to my parents. Without their support, this thesis would have

been impossible. I thank this country for its magnanimity in giving me a wonderful opportunity

that has provided me a memorable, valuable and an enriching experience. I thank my previous

colleagues Dr. Ramya Venkataraman, Dr. Prabhat Naredi, Dr. Venkata Pradeep Indirakanti, Dr.

Pramod Nimmatoori, Dr. Sudharshan Natarajan and Mr. Sridhar Ramanathan for their support

and encouragement during my stay at Pennstate.

1

Chapter 1

Introduction

1.1 Background

Carbon deposition, also known as coking or fouling, is common in systems when gaseous or

liquid hydrocarbons come in contact with metal surfaces at high temperatures and pressures.

Aviation gas turbines are not an exception to this problem. The thermal degradation of jet fuel is

more severe in the presence of dissolved oxygen [1-3]

. The fuel before combustion in the gas

turbine engine is also used as a coolant for thermal management of compressors, avionics,

hydraulics and environmental control [4]

. The amount of thermal energy that can be absorbed by

the fuel is limited by its thermal stability particularly in the high speed aircraft and is predicted to

increase by the aviation industry due to greater cooling requirements [5]

. The transport lines in the

fuel handling system are typically made up of Inconel alloys and stainless steel rich in nickel and

iron respectively. Fouling of Inconel and steel surfaces used in fuel injectors and other fuel

handling systems like servo mechanisms and heat exchangers increases the maintenance down

time of the aircraft. Therefore thermo-oxidative stability of jet fuel is a critical issue. Pressure,

temperature, fuel composition and the properties of metal surfaces play a key role in the

degradation of metal surfaces upon exposure to jet fuel [6,7]

. Carbonaceous solids can form on

metal surfaces from vapor phase and by metal catalysis under pyrolytic regime with different

morphologies [8]

and composition that are different from the carbonaceous mesophase that is

formed in the liquid phase [9]

.

1.1.1 Liquid phase autoxidation of hydrocarbons

Kinetic studies on autoxidation of organic compounds show the mechanism of

decomposition of hydrocarbons in the liquid phase [10-13]

. The oxidation of hydrocarbons in liquid

2

phase is a slow chain-branching reaction as opposed to combustion which proceeds through a

fast chain-branching reaction. Autoxidation of hydrocarbons is characterized by an initial

induction period, with an initially low rate of oxygen consumption followed by an increase in the

rate of oxygen consumption in subsequent stages. Molecular oxygen reacts with hydrocarbons by

a free radical chain reaction. The three main steps of the free radical chain reaction involved in

autoxidation are initiation, propagation and termination [14,15]

.

Initiation: The hydrocarbon molecules are activated by metals, heat, light or presence of

initiators like adventitious hydroperoxides in the medium to form a free radical (R·) by the

abstraction of hydrogen as shown in Reaction 1. As the chain length increases, there is a

reduction in bond dissociation energy for hydrogen abstraction from the alkyl radical. In liquid

phase autoxidation, the energy needed for abstraction of hydrogen from the hydrocarbon is

lowest for tertiary carbon and highest for primary carbon [16]

. In spite of being endothermic and

therefore thermodynamically and kinetically not favorable, Reaction 2 has been observed in the

autoxidation of indene [14]

.

RH R· + H· (1) Rate = [RH] [RI]1/2

RH + O2 R· + HO2· (2) Rate = [RH]3/2

[O2]1/2

RH + O2 + RH 2 R· + H2O2 (3) Rate = [RH]2[O2]

1/2

ROOH RO· + OH· (4)

2 ROOH RO· + RO2· + H2O (5)

As termolecular reactions are frequently possible in liquid phase reactions, Reaction 3 is most

likely when the bond dissociation energy for R-H is low. When the bond dissociation energy

3

increases, the enthalpy of Reaction 3 is higher than that of Reaction 2. Experimental observations

of higher rate of initiation in spite of higher C-H bond dissociation energy suggest that the

decomposition of adventitious hydroperoxides yields free radicals that can abstract hydrogen

atoms from the alkanes [16]

. Thermal decomposition of hydroperoxides form free radicals as

denoted by Reaction 4 and 5 that leads to chain reaction through autocatalysis. Unimolecular

decomposition (Reaction 4) occurs at lower concentrations of hydroperoxides. Dimerization and

the less endothermic bimolecular decomposition proceed at higher value of hydroperoxide

concentrations. Tertiary hydroperoxides are the most stable, whereas primary hydroperoxides are

the least stable compounds. Hydroperoxides decompose to form alcohols in acidic medium, and

carboxylic acids in a basic medium.

Propagation: Alkyl radical react with atmospheric oxygen rapidly to form a peroxy radicals

(Reaction 6). Subsequently, intermolecular hydrogen transfer (Reaction 7) from the hydrocarbon

to the peroxy radical, the rate limiting step in propagation of autoxidation, results in the

formation of hydroperoxides.

R· + O2 RO2 · (6)

RO2· + RH ROOH + R· (7)

The bond strength of the bond between hydrogen with oxygen in ROOH is estimated to be 90

kcal / mol [15]

. This is comparable to the bond energy of tertiary C-H bond in a saturated

hydrocarbon. The weak bond dissociation energy in the S-H, N-H and P-H bonds can also

provide hydrogen atoms to the peroxy radical to form hydroperoxides. The ability of hydrogen

abstraction of peroxy radical from a hydrocarbon also depends on the resonance stabilization of

alkyl radical and the availability of electrons in the carbon atom from which hydrogen has to be

4

abstracted. Alternatively, intramolecular hydrogen abstraction from the beta position of peroxy

radical forms a compound with a double bond. As the autoxidation progresses, the peroxy radical

would prefer to react with the oxygenated products rather than with the hydrocarbons as the

carbon hydrogen bond in the oxygenated product is weak [14]

.

Metal-Catalyzed Decomposition of Hydroperoxides: Metal cations promote the formation of

a complex with hydroperoxides and break them down into free radicals and ions by Reaction 8

and 9 that further propagate the oxidative degradation of hydrocarbons. Iron acts as an oxidizing

agent catalyzing Reaction 8 where as cobalt can catalyze both reactions.

ROOH + Mn+

RO· + M(n+1)=

+ OH- (8)

ROOH + M(n+1)+

RO2· + Mn+

+ H+ (9)

Termination of Hydroperoxides: At higher oxygen concentrations, chain termination occurs

by Reaction 10 due to the formation of peroxy radicals by the reaction between alkyl radicals and

oxygen.

RO2· + RO2· ROOR + O2 (10)

R· + R· Products (11)

1.1.2 Deposit Formation during Autoxidation of Hydrocarbons

1.1.2.1 Effects of Hydrocarbon Structure

The deposit formation tendency of the following normal paraffins decane, dodecane,

tetradecane, hexadecane, branched paraffins like 2,2,4,6,6 pentamethyl heptane, 2,3,4

trimethylpentane and aromatic compounds like tetrahydronaphthalene, 1-methyl naphthalene,

5

decahydronaphthalene and n-butyl cyclohexane and some of their blends were studied under

reduced pressure of 3 psia [17,18]

. The deposit forming tendency increased with temperature for

the normal paraffins and decreased with increasing carbon number in the temperature range

between 400 K and 500 K at a reduced pressure of 3 psia for the compounds investigated. This

experimental observation is in contrast to the fact that the oxidation rate increases with the

increasing number of secondary carbon atoms. For the same carbon number, the magnitude of

deposit formation was greater for a branched paraffin compared to normal paraffin both in pure

state and in binary blends with n-dodecane. This observation is supported by the fact that the

energy for hydrogen abstraction is the lowest for a hydrogen atom attached to a tertiary carbon

atom and the highest for that attached to a primary carbon. The deposit formation rate for pure 1-

methyl naphthalene was greater than that of pure n-decane above 120 ºC whereas the trend

reversed below this temperature The activation energy for deposit formation is higher for 1-

methyl naphthalene (21 kcal / mol) compared to that of n-decane (10 kcal / mol) which would

imply that addition of 1-methyl naphthalene would reduce the deposit forming tendency of the

binary blend. But the apparent activation energy for deposit formation changes with the

concentration of 1-methyl naphthalene in the blend. Among all the fractions of the blend at all

temperatures, there is an optimum concentration of 1-methyl naphthalene (approximately 10

wt% which varies slightly with temperature) in the binary blend that has the lowest deposit

formation rate. Below 120 ºC, addition of 1-methylnaphthalene inhibits the deposit forming

tendency of n-decane in the binary blend. Above this temperature, adding 1-methyl naphthalene

increases the deposit forming tendency of the blend. This suggests that the two entangled

variables, composition and temperature, play a complex role in the deposit forming tendency of

the mixture. The relationship between the deposit formation tendency of 10 % of aromatic or

6

naphthenic compounds in n-decane and number of benzylic hydrogen atoms in the aromatic or

naphthenic additive shown in Figure 1.1 implies that as the number of hydrogen atoms attached

to the carbon atom alpha to the single π (pi) electron system increases, the deposit formation

tendency of the blend decreases significantly and asymptotically relative to that of pure n-decane

suggesting that the blend containing naphthalene would produce more deposits and that

containing tetralin would produce the least amount of deposits. Compounds with hydrogen

atoms attached to the carbon atom between two single π electron system form a more stable

radical compared to that attached to the carbon atom alpha to the single π electron system.

Addition of olefins to n-decane has a deleterious effect on the formation of carbonaceous solid

deposits. Olefins may react with oxygen by addition reaction to form polyperoxides or by

hydrogen abstraction to form hydroperoxides. Olefins with conjugated unsaturation are more

susceptible to oxidation and show a greater tendency towards deposit formation

7

Figure 1.1. Effect of hydrocarbon structure in deposit formation17

1.1.2.2 Effects of Dissolved Oxygen

Experiments in the temperature range from 150 ºC to 690 ºC and at pressures as high as 69

atmospheres showed complex Arrhenius relationship for deposit formation of both oxygenated

and deoxygenated fuels with different levels of sulfur and nitrogen [19]

. This complexity is due to

the reduction in the concentration of deposit-forming precursors as the fuel changes from liquid

phase to supercritical phase with the increasing temperature. In the temperature range 150 º C –

560 º C, the deposit formation rate increases with temperature for both air saturated fuels and

deoxygenated fuels. However the amount of deposits formed decreases with fuel deoxygenation

8

especially in the liquid phase. Hazlett’s observation[9]

of the formation of 200 mg of

hydroperoxides and 0.01 to 0.1 mg deposits from one liter of n-dodecane agrees with that of

Taylor at least for low sulfur jet fuels[19]

. The oxygen content of deposits in the deoxygenated

fuel is significantly lower compared to that from air saturated fuels. Sulfur and nitrogen were not

in significant concentrations in the deposits. As the deposits had low molecular weight, it was

pointed out that the deposits formed in liquid phase depends upon the solvent characteristics

which changes when the oxygen atoms and some of the sulfur and nitrogen atoms get

incorporated in the deposits. The deposits formed in the gas phase were pointed out to be

dependent upon the molecular weight of the deposit precursor. The effect of pressure on deposit

formation is complex. The ability of deoxygenation to reduce the deposit formation depends

upon the trace level of sulfur compounds.

1.1.2.3 Effects of Sulfur Compounds

To elucidate the effect of sulfur in deposit formation, sulfur compounds were added to

deoxygenated JP-5 so that the dissolved oxygen and sulfur concentration are 0.3 ppm and 3000

ppm respectively and thermally stressed from 150 º C – 650 º C at 69 atmospheres[20]

. Among

the sulfide, disulfides, polysulfides and condensed thiophenes added, phenyl benzyl sulfide and

ditertiary butyl disulfide gave the highest amount of deposits although other compounds gave a

significant amount of deposits. An increase in the concentration of the two compounds caused a

non-linear increase in deposit formation. Condensed thiophenes inhibit the deposit formation

from jet fuels and do not increase deposit formation due to the better strength of aryl C-S bond.

Thiols, disulfides and dialkyl sulfides decompose on iron surfaces above 150 º C due to the weak

alkyl sulfur bond and S-S bond and diaryl sulfides decompose at 449 º C [20]

. Therefore they are

presumed to undergo surface catalysis to form deposits.

9

1.1.2.4 Effects of Nitrogen- and Oxygen-containing Compounds

Nitrogen compounds like 2,5-dimethyl pyrrole, indole, carbazole and quinoline were not

observed to have any effect on deposit formation when deoxygenated (0.1 ppm dissolved O2)

JP-5 with 100 ppm of each of these compounds was thermally stressed at 69 atmospheres in the

temperature range 371 ºC – 540 ºC [21]

. Cyclic ethers like furan, benzofuran and dibenzofuran,

carboxylic acid, alcohols and esters did not affect fuel stability whereas n-decanoic acid had a

deleterious effect in deposit formation. Irrespective of the molecular structure, 100 ppm of

hydroperoxides produced the same amount of deposits as 3000 ppm of sulfur compounds in jet

fuel. A steady rate of deposit formation was observed for air-saturated fuel without

hydroperoxides, whereas a sudden rise in the rate was observed for deoxygenated fuels with

external addition of hydroperoxides. This suggests that hydroperoxides are the precursors to

deposit formation. Naphthenic esters, alcohols, acids and ketones gave a relatively low amount

of deposits compared to their alkyl and aromatic counterparts. It was suggested that the products

formed from naphthenic compounds are more soluble than those formed from alkyl and aromatic

compounds. The 2, 5-dimethyl pyrrole and decanoic acid synergistically interacted under the

experimental conditions and produced deposits greater in amount than that produced by each of

the two compounds if present separately in JP-5.

1.1.2.5 Effects of Antioxidants

In addition to proposing the reactions that were discussed for autoxidation of hydrocarbons,

to take into account the effect of molecules which play the role of antioxidants Zabarnick has

considered the following reaction sequence [22]

.

RO2· + AH → ROOH + A· (12)

10

A· + O2 → AO2· (13)

AO2· + AO2· → Products (14)

A· + A· → Products (15)

AO2· + AH → AOOH + A· (16)

RO2· + A· → Termination (17)

The antioxidants (AH) intercept the peroxy radicals and prevent chain propagation by

Reaction 12. The antioxidant molecules are designed such that hydrogen abstraction from an

antioxidant molecule is relatively easier compared to that from the hydrocarbon molecule and the

resulting radical formed from the antioxidant is stable. Antioxidant radical (A·) would combine

with oxygen through Reaction 13 to form stable peroxy radicals due to their steric hindrance and

the molecular rearrangement. This property of antioxidant molecule would prevent chain

propagation reaction. The activation energy for the propagation Reaction 5 and the chain transfer

Reaction 12 determines the concentration of antioxidant to be added to the hydrocarbon to inhibit

autoxidation. It suggests that when any species that can get oxidized easily by losing a hydrogen

atom is added to the system the rate of autoxidation is reduced. This explains the reason for the

low oxidation rate of a jet fuel that is hydrotreated with less severity. The sulfur compounds act

as antioxidant species.

1.1.2.6 Synergism between Natural and Synthetic Antioxidants

Phenols, amines and thiols seem to acts as natural antioxidants and delay the autoxidation of

jet fuel [23,24]

. Studies show that synthetic antioxidant species provide readily abstractable

hydrogen atoms to the peroxy radical and intercept chain propagation as shown in Reaction 18.

11

The radical formed from the synthetic antioxidant species stabilizes itself through various

mesomeric forms as shown in Reaction 19 [25]

.

Natural antioxidants present in straight-run jet fuel donate abstractable hydrogen atoms to

the peroxy radical. But they also participate in the subsequent free radical reactions due to their

higher reactivity. Studies conducted to monitor the oxidation rate of Exxsol D110, a paraffin

based jet fuel that contains approximately 50% of paraffins, 50% cycloparaffins and less than 1%

of aromatics, after blending with straight-run fuel containing natural antioxidants show an

increase in the time taken for the consumption of 50% of dissolved oxygen. The same effect was

observed by adding butylated hydroxyl toluene (BHT) in air saturated Exxsol D110 at 185 ºC

[25]. Addition of 10 ppm of BHT and 5% of straight-run fuel to neat Exxsol D110 produced a

12

synergistic effect in reducing the rate of oxidation by increasing the time taken for the

consumption of dissolved oxygen. With the help of Reaction 20,

Anat· + AHhp → AHnat + Ahp· (20)

where Anat· and Ahp· are the antioxidant radical generated from the species present in straight-run

jet fuel and BHT respectively and AHnat and AHhp represent the antioxidant molecules present in

the straight-run and BHT respectively. This reaction prevents Anat· from propagating the chain

reaction.

1.1.2.7 Effects of Natural Antioxidants in the Autoxidation of Neat and Blended fuels

Autoxidation of jet fuels with different concentration of trace level heteroatoms has been

studied by measuring the time necessary for 50% of oxygen depletion. Straight-run fuels have

amines and phenols that act as antioxidants[26]

. These antioxidants reduce the rate of oxidation

for straight-run fuels. Hydrotreated fuels which do not have phenols, sulfur compounds and

amines oxidize very rapidly. An interesting observation of lower rate of dissolved oxygen

depletion for certain blends of jet fuels instigated the authors to find out the role of blending in

autoxidation. It was observed that the time necessary for depletion of 50% of initial dissolved

oxygen for certain blends, was far greater than that of the individual fuels. One of the

components of the blend is the hydrotreated fuel with lower amount of aromatics having lower

oxidative stability and the other component is a straight-run fuel with antioxidants and dissolved

metals having lower thermal stability. Blending the hydrotreated fuel with straight-run fuel

optimizes the antioxidant concentration and reduces the metal catalyzed autoxidation, thereby

improving the oxidative stability of the blend. However, the thermal stability of the blend does

not exceed that of the hydrotreated fuel. This process gives a lower amount of the carbon

13

deposits for short durations and a higher amount at longer durations compared to that for the pure

hydrotreated fuel. Blending the straight-run fuel with hydrotreated fuel also reduces the

concentration of antioxidants in the former and improves its thermal stability to some extent. A

fundamental understanding on the effect of blending in improving the oxidative stability and

thermal stability for fuels is incomplete.

1.1.2.8 Effects of Surface Catalysis on the Liquid Phase Autoxidation of Hydrocarbons

1.1.2.8a Effects of Metals

Taylor studied the catalytic effect of copper, pure titanium, titanium alloy, stainless steel

and aluminum on autoxidation of jet fuel at 204 º C with saturated air [27]

. The deposit formation

on metal strips was measured by the increase in weight of metal strips. The relative activity for

deposit formation was the highest for copper, moderate for titanium alloy and the lowest for pure

titanium, stainless steel and aluminum. The apparent activation energy for copper, stainless steel,

and titanium alloy was found to be 10 kcal / mole. Vanadium content in the titanium alloy

exhibited catalytic activity for deposit formation. The rate of deposit formation increased with

the increasing vanadium content. The catalytic activity of these metals was attributed to the

generation of free radicals on the surface of metals.

The degradation of fuel in the presence of oxygen is more severe compared to that in the

absence of oxygen[9]

. The rate of oxygen consumption during autoxidation on the surface of

aluminum was two-thirds of that in SS 304 and four-fifths of that in SS316 [28]

. Although, the

rate of formation of hydroperoxides is the lowest on aluminum and the highest on SS304, the

maximum concentration of the hydroperoxides is reached at the same temperature (260 ºC).

Beyond this temperature, the hydroperoxides decompose to form decanones and decanols, the

14

yield of the latter being higher than the former. As the reaction is carried out at higher

temperatures (500 ºC), alkanes, alkenes, carbon monoxide, hydrogen, methane, ethane and

ethylene are formed. The product distribution suggests the presence of three different regimes.

Autoxidation regime controls degradation until 260 ºC. Pyrolysis regime takes over beyond 480

ºC. Between 260 ºC and 480 ºC, there is an intermediate regime where the reaction mechanism

seems to be very complex[28]

. In the intermediate regime, apart from the decomposition of

hydroperoxides to alcohols and ketones, β-scission of alkoxy radicals formed during the

hydroperoxide decomposition forms aldehydes that subsequently decompose to carbon

monoxide and alkyl radicals. These radicals terminate to form lower alkanes and or undergo β-

scission to form 1-alkenes. The concentration of n-alkanes is greater than that of 1-alkenes in the

intermediate regime. This is due to the fact that β-scission of aldehyde forms lower alkane and

that of parent hydrocarbon forms a lower alkane and alkenes. At temperatures greater than 370

ºC, as the hydroperoxides have decomposed completely, there is shortage for hydroxyl radicals

and alkoxy radicals. This reduction decreases the rate of hydrogen abstraction and hence the

yield of lower alkanes and 1-akenes compared to that seen in the autoxidation regime.

QCM (quartz crystal microbalance) studies were conducted to measure carbon deposits

from jet fuel autoxidation at 140 ºC and 1atmosphere air due to the catalytic activity of the metal

for 15 hours suggest that a large amount of deposits was formed on platinum with lower amounts

of deposits obtained on the surface of gold and aluminum from hydrotreated fuels [29]

. It was

hypothesized that platinum catalyses the decomposition of hydroperoxides. A straight-run fuel

with 760 ppm of sulfur gave a larger amount of deposits on gold, aluminum and platinum.

15

1.1.2.8b Effect of Metal Oxides

Liquid phase autoxidation of n-hexadecane with copper oxide dispersed in the hydrocarbon

medium at 100 º C and 1 atmosphere initiated by tertiary butyl hydroperoxide showed that the

metal oxide exhibited a catalytic effect when its surface area is greater than 0.43 m2 per gram

catalyst in solution[30]

. For each catalyst surface area, there is a critical initial hydroperoxide

concentration below which the rate of oxidation is slow and above which the rate of oxidation of

n-hexadecane is fast. The critical initial hydroperoxide concentration is a strong function of

catalyst surface area. Experimental measurements of autoxidation rate using copper oxide

catalysts prepared by different methods revealed that Cu2+

and O2-

strongly accelerates

autoxidation, Cu2+

and OH- strongly inhibit autoxidation and Cu

+, O

- and O

2- strongly inhibit

autoxidation. The authors proposed a non-radical hydroperoxide decomposition mechanism on

Cu2+

and O2-

surfaces and initial free radical decomposition followed by subsequent non-radical

decomposition due to the conversion of Cu2+

to Cu+ during the decomposition of hydroperoxides

[30]. Silcosteel coating was produced by Restek Corporation by chemical vapor deposition of a

proprietary silica based layer on stainless steel surfaces. Studies conducted with Jet-A on

silicosteel and stainless steel tubes show that the rate of consumption of dissolved oxygen and

hence the amount of solid carbonaceous deposits are lower for treated tubes relative to that for

untreated tubes [31]

. This was attributed to the absence of active sites on the treated tubes that

might participate in the decomposition of hydroperoxides. The combination of surface coating

and an additive package shows a synergistic reduction in the amount of deposits. Even though

thick solid carbonaceous deposits have reduced the activity of the surface towards carbon

deposition during the long duration of thermal stressing, surface treatment has a more

dominating effect in reducing the amount of deposits.

16

1.1.2.9 Autoxidation of Jet Fuels and Deposit Formation

The ability of a fuel to resist the formation of deposits may be defined as thermal stability of

the fuel. The oxidative stability of the fuel is defined as its ability to resist oxidation.

Hydrotreated fuels have lower oxidative stability and higher thermal stability compared to

straight-run fuels. Comparison of the Arrhenius plot of autoxidation products from thermal

stressing of air saturated n-dodecane[28]

with that of the deposition rate[9]

suggested that the

hydroperoxide decomposition, autoxidation product formation and deposit formation ramp up at

262 ºC. The H/C ratio of the deposits present in the actual turbine engine indicates that the

deposits are substantially made up of aromatic and heteroaromatic compounds[9]

. Hazlett

proposed that hydroperoxides are the precursors to deposit formation. A comprehensive

mechanism was proposed by Kauffmann to address the formation of deposits on both

hydrotreated and straight-run jet fuels[23]

. In addition to the reactions 2, 4, 5 and 12 the following

reactions are proposed for deposit formation.

RCHO + RCOOH → Polymers (19)

ROOH + ( RSR / RSSR ) → Phenols + acidic sulfur oxide + RCOR (20)

Acidic sulfur oxide + basic nitrogen compounds → S•N Compounds (21)

S•N Compounds + Phenol radicals → Bulk particles (22)

Acidic Sulfur oxide + Metal surface → Initial deposition (23)

Initial Deposition + Bulk Particles → Surface deposition (24)

Addition of reactive sulfur compounds like phenyl benzyl sulfide and diphenyl disulfide to

hydrotreated Jet A1 and straight run fuel Jet A produced significant amount of phenols and fewer

17

amount of hydroperoxides during autoxidation at 160 ºC. Addition of mercaptans inhibited

oxidation, produced fewer phenols and less hydroperoxides. Sulfoxides and sulfones which are

supposed to be the autoxidation products gave rise to the formation of significant amount of

hydroperoxides. XPS and infrared spectroscopy of the particles from autoxidation of

hydrotreated fuel suggested the presence of relatively large amount of carbon to oxygen double

bonds formed from aldehydes and acids, whereas those from straight-run Jet A has substantial

amount of sulfur oxides, nitrogen and carbon to oxygen single bonds. Depth profiling of initial

surface deposits on the metals obtained from stressing of hydrotreated fuel Jet A-1 at 210 ºC for

10 minutes showed 9-16% sulfur and less than 1% nitrogen with sulfides and disulfide additives

and less than 1% of sulfur and nitrogen with sulfones and sulfoxides additives. A similar analysis

on deposits obtained from stressing of straight-run jet fuel at 210 ºC for 10 minutes in the

presence of iron oxide, calcium oxide, aluminum oxide and silica shows that acid neutralizing

oxide compounds reduce the amount of deposits on metal surfaces. FTIR, gas and liquid

chromatography measurements of alcohol, ketones and other oxidative products suggests that the

hydrotreated fuel undergoes oxidation to a greater extent[22]

. These observations support the

mechanism proposed[23]

.

The experimental observations made by Hazlett[9,28]

between 260 ºC and 430 ºC were

contrasting to that by Taylor[17]

at 121 ºC. Qualitative information from FTIR about the deposits

formed at 450 ºC from jet fuel revealed the presence of carbon-oxygen single and double bonds

and aliphatic and aromatic carbon-hydrogen bonds. Mass spectrometry results of the deposits

revealed the presence of napthalenes, aromatics with one oxygen atom, aromatics with one

nitrogen atom, aromatics with sulfur atom, aromatics with nitrogen and oxygen atoms and

18

aromatics with two oxygen atoms[9]

. Hazlett proposed a mechanism for the formation of deposits

given in Figure 1.2.

Figure 1.2. Chemical and physical process in thermal deposition9

The mechanism proposes that autoxidation of jet fuel involves all the species that produce

soluble low molecular weight products containing 8%-10% of oxygen, nitrogen and sulfur

atoms. Further autoxidation leads to the formation of products with molecular weight as high as

200 to 600 Daltons. These products being insoluble in jet fuel precipitate out as deposits. The

deposit precursors agglomerates in the liquid phase, gets anchored to the surface and fuse to form

varnish. Sulfur and nitrogen compounds are easily oxidized relative to the n-alkanes. Oxidation

of sulfur by peroxides rather than dissolved oxygen would yield a variety of products. One

theory proposes that incorporation of heteroatoms into soluble products during autoxidation

19

increases their polarity and reduces the solubility in the non-polar jet fuel. Spectroscopic

characterization of the deposits supports this hypothesis. As opposed to this, another theory

proposes that deposits are formed by the polymerization of olefins and pyroles to form high

molecular weight products which precipitate as deposits[32]

. Molecular weight determination of

the deposits indicates less polymerization and this theory is not well supported. The mechanism

proposed by Hazlett has been extended to account for the formation of thermal-oxidative deposit

in jet fuels in scheme 1[33]

. In scheme 1, the peroxy radicals produced during autoxidation of jet

fuels are intercepted by the natural antioxidant molecules present in the fuel to produce

hydroperoxides. The phenoxy radicals formed during the decomposition of natural antioxidant

reacts with dissolved oxygen to form keto- peroxy radicals. The keto-peroxy radicals dimerise to

form tetroxide molecules which decompose to form hydroquinone type compounds and quinone

type compounds evolving molecular oxygen. Quinone being an electrophile undergoes

electrophilic aromatic substitution with electron rich carbazole to form a soluble macromolecular

oxidatively reactive species (SMORS). This species reacts with hydroquinone species class

molecule to form a compound which would oxidize further to precipitate a high molecular

weight species. These particles agglomerate to form deposits[33]

. Silcosteel surfaces reduce

deposit formation during oxidative degradation of jet fuels34

suggesting a role for surface

catalysis. To take surface effect into account, a Scheme 2 is proposed where an aryl thiol species

adsorbs on the metallic surface forming a sulfide which would further react with the products of

oxidative degradation to form surface deposits. This proposed mechanism has accounted for

some experimental observations [33]

. Thus the gap between autoxidation of jet fuel and the

formation of deposits is well bridged by the mechanisms proposed.

20

1.1.3 Factors Affecting Deposits under Pyrolytic conditions

Fuel composition, presence of sulfur compounds and catalytic nature of substrate surface

are the important factors that contribute to carbon deposition in pyrolytic regime. Free radical

reactions initiate the formation of alkenes that undergo cyclization to form alkylcyclohexenes

which upon subsequent dehydrogenation and condensation produces polycyclic aromatic

hydrocarbons[35]

. The pyrolytic decomposition of long chain alkanes can be retarded by the

presence of hydrogen donors like tetralin and decalin. At high temperatures, these donors

undergo degradation by saturated ring rupture and dehydrogenation that would ultimately lead to

the formation of naphthalene[36]

.

The influence of sulfur on carbon deposition is very complex. Addition of organic sulfides

in the reactant stream appeared to suppress carbon deposition during ethane pyrolysis over Fe/Ni

(1:4) catalyst particles[37]

. Characterization of solid deposits from thermal stressing of n-

dodecane in the presence of thiophene, 3-methyl benzothiophene and benzyl phenyl sulfide on

metal substrates suggests two contrasting observations[38]

. In the presence of Fe and Ni, all the

three organic sulfur compounds inhibited carbon deposition ascribed to the blockage of active

sites by sulfur poisoning, whereas in the presence of inconel 718, benzyl phenyl sulfide

promoted carbon deposition which is ascribed to the disruption of alloy surfaces due to the

formation of iron and nickel metal sulfides. The nature of metal sulfide formed during the

interaction of reactants with the catalytic surface affects the rate of carbon deposition[39]

.

Formation of iron sulfide in stainless steel and iron surface passivates the metal surfaces and

reduces the rate of carbon deposition whereas the decomposition of nickel sulfide followed by

surface disruption increases the rate of deposition from propane under pyrolytic conditions.

Exposure of the active metal due to the disruption of metal sulfides and increase in surface area

21

due to the formation of metal sulfides can increase the amount of carbon deposits. This

proposition invoked by Turner etal[39]

is supported by the experimental observation where the

amount of carbon deposits increase with sulfur concentration in Jet-A due to higher surface area

offered by metal sulfides for carbon deposition[40]

. On the other hand the formation of lower

amount of carbonaceous deposits on catalytically active surfaces from Jet-A relative to the

inactive substrate supports the proposition that the sulfur compounds in Jet-A block the active

sites on the catalytic substrates[41]

.

1.1.4 Solid Deposit Formation from Jet Fuel Obtained under Pyrolytic Conditions

Solid deposits with different morphologies were obtained on different superalloy surfaces

under the same conditions of thermal stressing of jet fuel. The formation of carbon

nanostructures from jet fuel on the surface of inconel 600 and inconel 718 at 35 atmospheres and

500 ºC was first reported[8,42]

. The temperature and pressure conditions of the formation of these

nanostructures differ from the typical low pressure and high temperature conditions needed for

their growth. Plate-like structures were formed on the surface of Havar[43]

. Filaments and

amorphous carbon deposits were formed on Waspalloy surfaces[43]

.

1.1.4.1 Carbonaceous Mesophase

There is comprehensive investigation on the kinetics of formation, development of

morphology and chemical composition of mesophase carbon from various organic precursors

and their mixtures[44-49]

. Carbonization and graphitization of organic materials are controlled by

the chemical structure of starting organic material[46]

. Carbonaceous mesophase forms in the

temperature range between 350 ºC and 500 ºC [50]

. The mechanism of formation of carbonaceous

mesophase from the parent isotropic pitch was discussed in detail in the literature[44,45,50,51]

. Non-

22

aromatic structures are aromatized during carbonization. Subsequently it gives rise to the

formation of free radicals by C-H, C-C bond cleavage, molecular rearrangement, thermal

polymerization and aromatic condensation followed by the elimination of side chain and

hydrogen[46]

. During carbonization, the concentration of large molecules increases in liquid

phase. At a critical concentration, there is a transition from liquid phase to the formation of a

liquid crystal phase (mesophase) characterized by the reduction in the mobility of molecules.

Mesophase is a chemically reactive phase such that an increase in the carbonization temperature

results in polymerization reaction of many polyaromatics hydrocarbons leading to the formation

of semi coke. The large molecules are held together by van der Walls forces[50[

. Oxidation and

presence of sulfur at high concentrations inhibits mesophase development[46]

. The growth of

mesophase depends upon the time and carbonization temperature of the system and is

independent of the nature of mixing of the growth medium. High viscosity of the medium gives

rise to smaller areas of anisotropic structures called mosaics and low viscosity gives rise to

domains. High carbonization temperature and lower time have similar effect on the development

of larger areas of anisotropic structure of the mesophase as would low temperature and long

duration of time. Anisotropic botryoidal spheres with different sizes form at different

carbonization temperatures depending upon the reactivity of the starting precursor. The

development of botryoidal spheres is a balance between the rate of growth of spheres, increase in

viscosity of mesophase due to polymerization and decrease in viscosity due to higher heat

treatment temperature.

The formation of carbonaceous mesophase was also observed in deposits formed on the

aircraft fuel lines[52]

. Globular particles formed on the afterburner fuel are characterized by the

formation of isotropic matrix with planar molecules arranged approximately parallel to each

23

other without any stacking sequence. This structure confirms the fact that the formation of

carbonaceous mesophase is controlled by the rate of liquid phase reactions. The paraffin

molecules undergo cracking, cyclization, and polymerization to form large polyaromatics

compounds. Once these planar molecules are formed, the mesophase formation requires higher

residence time and high temperatures. The formation of mesophase indicates that the fuel must

have been stagnant in the afterburner fuel line for a long time. This could be possible as the

afterburner line is used only when high thrust is needed to the aircraft during rake-off and high

speed maneuvering and the lines are not used during cruising. The formation of mesophase is not

expected in deposits from flow reactors where the residence time is very low.

1.1.4.2 Filamentous Carbon

The formation of filamentous carbon from light hydrocarbons on the surface of iron and

nickel and their properties has been studied extensively[53-55]

. Filamentous carbon forms in the

temperature range between 400 ˚C and 1000 ˚C from gaseous species on the transition metal

surfaces such as Iron and Nickel. The hydrocarbon gases undergo adsorption, dissociation on the

surface of metal, dissolution and saturation of the metal by carbon. This is followed by the

precipitation of carbon from the other side of the metal that results in the formation of filaments.

The nature and composition of the metal surface is one of the many factors that affect the

morphology of carbon deposits. The alloy surfaces used in the fuel handling system are

predominantly rich in iron and nickel. These metals are known to exhibit catalytic activity

towards hydrocarbons that would lead to the formation of filamentous carbon[56-58]

. The

formation of filamentous carbon from jet fuel on Inconel 600 has been reported[8]

. Flow reactor

studies on a various metal substrates and characterization by temperature programmed oxidation

showed deposits that vary widely from structurally disordered carbon to highly ordered carbon

24

deposits[43]

. Presence of aluminum, titanium, niobium and tantalum in the Inconel 750-X appears

to suppress the catalytic activity of nickel towards carbon deposition. Fe/Ni ratio in each

substrate also appears to play an important role in deposit formation[41]

. Characterization of

deposits on inert surfaces like glass-lined tubing and Silcosteel shows the absence of highly

ordered carbon due to the absence of catalytic activity toward carbon deposition[59]

.

1.1.4.3 Spherulitic Deposits

The mechanism of formation of carbon in the gas phase is documented in carbon

literature[60-62]

.

The pyrolysis of hydrocarbons in the gas phase produces a number of

intermediate species. The complexity of the intermediate species varies with molecular weight

and C/H ratio. The lesser the complexity, the lower the melting point and viscosity of the

species. Collision and condensation of intermediate species result in the formation of droplets

with sufficient complexity. This is followed by nucleation and growth that take place

simultaneously. Droplets with sufficient fluidity and low residence time in the gas phase tend to

form a continuous film upon collision with the solid substrate. The rigid droplets with larger

residence time coalesce in the gas phase with the fluid matrix. The strength of the bond between

the droplet and the matrix is different from that within either one of the phase. Higher residence

time of the droplets makes the skin rigid due to higher degree of polymerization and prevents

further coalescence. These droplets retain the spherical shape and do not wet the surface of the

substrate upon collision. The degree of crystallite orientation depends upon the pyrolysis

temperature. Spherulitic carbon deposits have been observed from Jet-A under pyrolytic

conditions on SS316[63]

and autoxidative conditions[31]

. These structural order of these deposits

depend on the temperature at which the respective studies were conducted.

25

1.1.4.4 Pyrolytic Carbon

Hydrocarbons undergo pyrolysis in gas phase and give rise to the formation of solid

carbonaceous deposits either by homogeneous nucleation resulting in the formation of soot or

heterogeneous nucleation resulting in the formation of surface nucleated deposits. Low

molecular weight hydrocarbons undergo dehydrogenation, nucleation and growth to form

polycyclic aromatic hydrocarbons (PAH) which are considered as the precursors to the formation

of soot[64]

. These precursors appear to undergo condensation further by Dies-Alder reaction or

free radical reaction[65]

. Beyond 1300 °C, the formation of soot is involves the loss of hydrogen

atom from an aromatic ring followed by the successive addition of acetylene to close the

aromatic ring[66]

. The precursors to the formation of surface nucleated carbon is not clear.

Defects on the surface of the substrate assist in the initiation of surface nucleated deposits.

Surface nucleation of the precursors from the gas phase on these defects results in the formation

of growth cones. When all the growth cones originate at the surface, they are known as

singularly nucleated deposits. When the nuclei form throughout the deposits, it results in

continuous or regenerative nucleation[67]

. Pyrolytic carbon was observed in the burner line and

the afterburner line of an aircraft engine[52]

. These deposits are not expected to be formed in this

study where the thermal stressing temperature is low.

1.1.4.5 Metal sulfides

The reaction between organosulfur compounds and metals rich in iron and steel is expected

to result in the formation of the respective metal sulfides. Iron sulfide exists in the form of

pyrrhotite (Fe1-xS), trolite (FeS), mackinawite (Fe1+xS) and pyrites (FeS2). The formation of

pyrrhotites due to the reaction between organic sulfur compounds and metal surfaces rich in iron

26

and nickel is very common in petroleum refineries at temperatures above 250 ˚C [68]

. They also

form when sulfur compounds in Jet-A react with iron and nickel rich alloys under pyrolytic

conditions[40]

. These filaments were found to be 100 nm in diameter and several microns in

length. Pyrrhotites exist in the form of faceted crystallites and filaments that is also known as

nanorods. These sulfides belong to chalcogenides and of more interest due to its nanostructure.

The value of x in pyrrhotites varies between 0 and 0.125. Iron sulfide being a magnetic material

and having application in information storage, sensing and imaging has recently drawn attention

toward its synthesis in the form of various nanostructures such as nanowire, nanorods, nanosheet

and nanoplates by vapor, vapor-liquid-solid and solution-liquid-solid methods[69,70]

. Iron sulfides

were synthesized by boiling a mixture of iron oleate, alkylthiols and organic solvents like

tetradecane, 1-hexadecene and 1-octadecene at the respective boiling points in the temperature

range between 265 ˚C and 340 ˚C [71]

. It was found that iron sulfide formed from elemental sulfur

did not have any distinctive morphology whereas the ones from alkylthiols formed hexagonal

nanoplates and nanorods.

Figure 1.3. Schematic showing the growth mechanism of Fe-S nanostructure (adapted and

modified)71

27

The mechanism in the formation of nanorods or nanoplates shown in Figure 1.3 involves the

formation of hexagonal discs in the initial stages, oriented assembly of these discs in the form of

plate-like nanoflower in the intermediate stage by edge to edge fusion and growth of nanoflowers

along the c-axis to form nanorods. The edge to edge fusion takes place at temperatures between

320 ˚C and 340 ˚C that would result in nanorods or nanoplates. According to Bravais’ law,

crystal faces with low index and large interplanar spacing forms first in the process of crystal

growth. Therefore the law predicts that hexagonal structures prefer to grow along [100], [110]

and [00x] directions[72]

. The axis of the nanorods and their growth direction observed during

APCVD of pyrrhotites follows the Bravais’ law[72]

. Examination of filaments by transmission

electron microscopy (TEM) in the solid deposits formed on Inconel 600 from Jet-A under

pyrolytic conditions showed the formation of heazlewoodite (Ni3S2) and pentlandite (FeNi)9S8

and pyrrhotites on SS316. The filamentous or nanorod like morphology of heazlewoodite has not

been reported elsewhere.

1.2 Objectives of the Thesis

The principal objectives of this study are:

To examine the effect of different alloy surfaces on the formation of solid deposits from

Jet-A and examine the nature of solid deposits on metal surfaces in the intermediate regime

where both autoxidation and pyrolysis play an important role in deposit formation.

To evaluate a non-line-of-sight deposition process MOCVD (metalorganic chemical

vapor deposition) to coat tubes of varying diameters.

To evaluate the performance of coating candidates in inhibiting solid deposit formation

from jet fuel.

28

1.3 Organization of the Thesis

In this thesis, the results are presented in four different chapters. An abstract or

background, experimental, results, discussions and the reference sections are included for each

chapter. The figures and tables can be found after the references for each chapter. Chapter 2

focuses on the analysis of carbonaceous deposits formed from Jet-A on seven different metal

substrates to examine the effect of metal surfaces on the formation of solid deposits in the

intermediate regime. Chapter 3 reviews the various methods typically used to deposit

environmental barrier coatings, the details of the system configured for surface coatings used in

this study, and the evaluation of the MOCVD process to coat tubes of varying diameters and

identification of the correct precursor that would produce coatings of desired morphology

suitable for the inhibition of solid deposit formation from Jet-A. In Chapter 4, the effectiveness

of five coating candidates is investigated in inhibiting solid deposit formation. Chapter 5

evaluates the nature of carbon incorporated in the coating by various characterization methods

and the effect of carbon in zirconia coating in inhibiting carbon deposition from jet-A.

29

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30

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Fuel in a flow reactor. Industrial and Engineering Chemistry Research 2001, 40, 589-595.

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development. Part 2. Co-carbonaization of aromatic and organic dye compounds, and influence

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45. Marsh, H.; Foster, J. M.; Hermon, G.; llay, M.; Melvin, J. N., Carbonization and liquid

crystal mesophase development. Part 3. Co-carbonization of aromatic and heterocyclic

compounds containing oxygen, nitrogen and sulfur. Fuel 1973, 52 (4), 243-252.

46. Lewis, J. C., Chemistry of Carbonization. Carbon 1982, 20 (6), 519-529.

47. Honda, H., Carbonaceous Mesophase: History and Prospects. Carbon 1988, 26 (2), 139-156.

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properties of the products-II. Discontinuous reactor. Carbon 1992, 31 (1), 1-8.

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34

Chapter 2

Analysis of Carbonaceous Solid Deposits from Thermal Oxidative Stressing of Jet-A Fuel

on Iron and Nickel-based Alloy Surfaces

2.1 Abstract

Thermal stressing of Jet-A was conducted in a flow reactor on iron and nickel-based metal

surfaces at a fuel flow rate of 1 mL/min for 5 hours at a wall temperature of 350 °C and 3.5 MPa

(500 psig) so that both decomposition of oxidation products from liquid phase autoxidation and

pyrolysis contribute to the formation of carbonaceous deposits. The deposits produced were

characterized by field emission scanning electron microscopy (FESEM) and temperature

programmed oxidation (TPO). The effect of metal surface on deposit formation increases in the

following order AISI316 < AISI 321 ~ AISI 304 < Inconel 600 < AISI 347 < Inconel 718 <

FecrAlloy < Inconel-750X. The variation in the activity of the metal substrates is attributed to

their reaction with reactive sulfur compounds in the fuel and interaction of oxygenated

intermediates formed by autoxidation during thermal stressing.

2.2 Introduction

The formation of carbonaceous deposits from jet fuel on the metal surfaces in the fuel

systems before combustion is of major concern for the operation of aircraft engines as it can plug

the filters and accumulate on valves, flow lines and fuel injector [1,2]

. Temperature is one of the

important parameters affecting the rate of fuel degradation. Autoxidation of jet fuel which causes

the formation of oxygenated products is predominant at temperatures less than 260 °C [3]

. These

products decompose between 290 °C and 350 °C beyond which pyrolysis is significant. Fuel

composition strongly affects the thermal and oxidative degradation of the fuel and its deposit

35

forming tendency [4]

. Below 260 °C, the deposit forming tendency from jet fuels saturated with

air was significantly affected by the nature of exposed metal surfaces [5]

. It was observed that

metals containing copper and vanadium were most active toward carbon deposition.

Characterization of deposits formed at 260 ºC from Jet-A in the presence of excess air on metal

and metal oxide substrates showed similarities in its morphology and chemical composition with

soot [6]

. The morphology of deposits indicated negligible role of substrates in carbon deposition.

The presence of oxygen-containing functional groups and absence of sulfur and nitrogen in the

deposits was notable [6]

. Pyrolytic degradation studies conducted with commercial aviation fuel

Jet-A and military jet fuel JP-8 at 500 °C suggested that the amount of deposits formed is

strongly influenced by the nature of the metal substrate, the fuel composition and of any sulfur

compounds present [7]

. Investigation on the effect of substrate on deposit formation from JP-8 at

the same temperature has shown the presence of carbon with different level of structural order

ranging from amorphous to crystalline phases [2,8]

. It was also suggested that the presence of

minor components like Ti, Al, Nb and Ta in the alloy decreases the catalytic activity of iron and

nickel by reducing the solubility of carbon in the base metals and suppressing carbon deposition

through stabilization of alloy surfaces making the removal of metal particles from the surface

more difficult [2]

.

The interaction of sulfur compounds with metal surfaces is very complex [9]

. In some cases,

it was suggested that they can passivate the metal surface by forming sulfides and block the

active sites [7]

. On the other hand, it was suggested that sulfur compounds in jet fuel activate the

metal surface for carbon deposition by forming metal sulfides under pyrolytic conditions and so

increasing the surface area available for carbon deposition [10]

. Characterization of the deposits

obtained during thermal-oxidative degradation of n-hexadecane at 160 °C showed the formation

36

of aromatic solids in the fluid phase [11]

. Experiments primarily focused on the product formation

from the decomposition of aerated dodecane at 800 psig in the liquid phase in SS304, SS316 and

aluminum tubes suggested that formation of oxygenated products from hydroperoxide

decomposition and hydroperoxide initiated pyrolysis are predominant in the temperature range

between 282 °C – 400 °C [12]

. Although there were some differences in the product distribution,

the type of metal surfaces did not appear to control the type or amount of product formation.

Studies conducted with Jet-A under thermal-oxidative conditions where the fuel exit temperature

is less than 350 °C shows that surface reactions affect the carbonaceous solid deposit formation

[13]. The objective of this study is to investigate the effect of various metal alloys AISI304, AISI

316, AISI 347, AISI 321, FeCrAlloy, Inconel 600, Inconel 718 and Inconel 750-X on carbon

deposition from Jet-A at a wall temperature of 350 °C and a reactor pressure of 500 psig.

2.3 Experimental Section

2.3.1 Thermal Stressing Experiments

The elemental composition of eight foil substrates in weight percentage used in this study

is given in Table 2.1. All the substrates are washed in hexane and dried in argon for an hour

before the experiment. The experimental setup for thermal stressing of Jet-A is shown in Figure

2.1 [14]

. The details of the thermal stressing reactor are described elsewhere [1]

. The stressing

experiment is conducted in a 6.35 mm diameter (1/4-in o.d.) glass-lined stainless steel reactor

that is 20 cm long. The substrate is inserted at the bottom of the isothermal glass-lined stainless

steel reactor. The reactor with the foil is heated in the presence of argon at a reactor pressure of

3.5 MPa (500 psig) to 350 °C with the help of a block heater to maintain isothermal conditions

along the length of the reactor and maintained at that temperature for 4 hours to obtain thermal

37

equilibrium. Ultra zero air is bubbled into the Jet-A reservoir so that it is saturated with dissolved

oxygen during the course of the experiment. The fuel is pumped into the system at 500 psig. It

enters the preheating line of 3.175 mm diameter (1/8-in o.d.) and 2 m in length. The residence

time of the fuel in the preheating line is 6.3 minutes. It is preheated to 260 °C to initiate the

autoxidation before entering the reactor. The fuel flow rate, reactor wall temperature and the

pressure are maintained at 1 mL/min, 350 °C and 500 psig for 5 hours. The residence time of the

fuel in the reactor is 1.4 minutes. The fuel is maintained in the fluid phase during the course of

experiment. At the end of the experiment, the residual fuel in the reactor was removed by

purging it with argon.

2.3.2 Characterization of Carbon Deposits

Under pyrolytic conditions, maximum deposition was obtained between 10 and 15 cm from

the top of the reactor [1]

. Preliminary experiments were conducted with AISI 316 and AISI 304

foils to study the variation in deposit formation as a function of reactor length. The length of the

foils used in this experiment was 10 cm. With respect to the glass-lined reactor described

elsewhere [1]

, it is located between 7.5 cm and 17.5 cm from the top of the reactor. TPO

conducted on the two sections of the same substrate material, each 5 cm long, showed the same

amount of deposits. The nature of the TPO curves corresponding to each substrate is discussed in

the next section. The TPO curves from each of the two sections of the same substrate were

similar. Therefore, for the TPO of the other substrates, the portion of the foil located between 10

and 15 cm of the top of the reactor was chosen for analysis. The samples are dried under vacuum

at 110 °C for 2 hours. The morphology of the deposits was examined using a field emission

scanning electron microscope (FESEM) JEOL 6700F. X-ray diffraction was performed in the

gracing incidence mode to identify the various phases of metal sulfides if any in the deposits in

38

the PANalytical X’Pert Pro MPD instrument operated at 45 kV/40 mA and scanned at 0.02 º/s.

The amount of solid carbonaceous deposits formed on each substrate after 5 hours of thermal

stressing experiment is measured by temperature programmed oxidation (TPO) in a RC412

Multiphase Carbon Analyzer. During TPO, the sample is loaded in a quartz boat and heated from

100 °C to 900 °C in the presence of ultra high purity oxygen at a ramp rate of 30 °C/min and

held at 900 °C for 5 minutes. The carbon dioxide produced is measured in an IR cell. Any CO

produced during oxidation is converted to CO2 in the presence of a copper oxide catalyst. The

peak positions relate to the oxidation reactivity and thus depend on the structure of solid coke

deposits. The ramp rate during the TPO experiment may influence the position of peaks during

the oxidation of carbonaceous deposits [15]

. However, staged TPO experiments show that the

structure of the deposits does not change when the sample is heated in the above mentioned

program sequence in presence of UHP oxygen [16]

. The individual peak positions and peak

intensities are reproducible. The total amount of solid carbonaceous deposits obtained on each

substrate is reproducible to within 10% of the deposit mass.

2.4 Results and Discussion

The analysis of Jet-A by gas chromatograph-mass spectrometry (GC-MS) and GC with

pulsed flame photometric detector (GC-PFPD) for the hydrocarbon composition, and sulfur

compounds are shown in the chromatograms in Figure 2.2a and 2.2b respectively. The

concentration of sulfur compounds in Jet-A was found by elemental analysis to be 1160 ppm by

weight. By comparison with standards, some of the peaks in the chromatogram were identified as

dimethyl and trimethyl benzothiophenes as shown in Figure 2b. Characterization of sulfur

compounds in aviation fuels by atomic emission detector (GC-AED) has shown, in the order of

increasing retention times the presence of thiols, sulfides, disulfides (classified as reactive sulfur

39

species) and methyl-substituted thiophenes and benzothiophenes (classified as non-reactive

sulfur species). The classification of reactivity of the sulfur compounds is based on their

tendency to undergo hydrodesulfurization [17]

. Therefore, the unidentified peaks observed in the

chromatogram with shorter retention times may correspond to reactive sulfur species such as

sulfides and disulfides.

2.4.1 Amount of Solid Carbon Deposited on Different Metal Substrates

Based on the average amount of deposits obtained from three experiments on each

substrate, the variation in the tendency of each substrate for deposit formation from Jet-A is

shown in Figure 2.3. The reproducibility data for the TPO profile of solid carbonaceous deposits

on all the substrates are shown in Appendix A. The amount of carbon deposits formed on each

substrate is more than that needed to form a monolayer as seen from the calculations shown in

Appendix B. The amount of carbon deposited is lowest for AISI 316 and highest for Inconel

750-X. Among the stainless steel substrates, AISI 347 gave the highest amount of deposits. The

elemental composition of AISI 304 and AISI 347 shown in Table 2.1 is very close to one another

except for the presence of Niobium in AISI 347. Thermal stressing experiments with JP-8 on

niobium foils under pyrolytic conditions suggest that niobium is catalytically inactive for carbon

deposition [8]

. Similarly, the amount of carbon deposits on AISI 321 is close to that on AISI 304.

The composition of major elements and some minor elements is similar to one another except for

the presence of Titanium in AISI 321. Even though, titanium is known to suppress carbon

deposition under pyrolytic conditions, its presence does not make a difference in carbon

deposition from Jet-A at 350 °C as compared to other stainless steel foils. It should be pointed

out that the iron content of AISI 304 and AISI 321 (69% and 68%, respectively) is higher than

that of AISI 316 (64%) which deposits less carbon than those on AISI 304 and AISI 321.

40

Comparison of the elemental composition of Inconel 600 and Inconel 750-X shows the

presence of minor elements titanium, niobium and aluminum in Inconel 750-X in small

percentages which otherwise has a similar composition of major elements. Inconel 750-X gives

greater amount of deposits compared to Inconel 600. Under pyrolytic conditions, it was observed

that the presence of minor elements like Nb, Al and Ti in the metals, which are added for

precipitation strengthening [2]

, appeared to suppress carbon deposition [2]

. The ability of these

elements to suppress carbon deposition in the metals was attributed to the formation of

passivating layers that prevent the access of reactive species, formed during thermal

decomposition of hydrocarbons [2]

, to the base metals iron and nickel which are known for their

catalytic activity toward carbon deposition [18]

. The difference in the amount of deposits was

attributed to the surface composition of the alloy substrates which influences their catalytic

activity for deposit formation at 500 °C. The same explanation does not appear to support the

results obtained at 350 °C. As the temperature in these experiments is substantially lower than

that of pyrolytic conditions, the metals might exhibit a lower degree of catalytic activity toward

deposit formation from pyrolytic decomposition of hydrocarbons. In the intermediate regime, the

formation and decomposition of oxygenated products both in the fluid phase and on substrate

surface as well as the formation of deposit forming precursors in the fluid phase may contribute

to coke deposit formation.

2.4.2 TPO and FESEM Analysis of Deposits on Various Substrates

The peaks corresponding to the evolution of CO2 at low temperatures are due to the high

oxidation reactivity of hydrogen-rich carbon, or solid carbon that is structurally less ordered. The

CO2 peaks evolving at relatively high temperatures are due to the presence of hydrogen-lean

carbon that is structurally more ordered. The sulfur compounds present in Jet-A are expected to

41

react with the metal substrates to produce metal sulfides. An Fe-Ni-S ternary phase diagram at

400 ºC [19]

was used to predict the phases of various metal sulfide structures observed in this

study on various substrates. At 350 ºC, it is assumed that all the sulfur in the fuel is converted to

metal sulfides during the five hour duration of the experiment. This assumption is used to

calculate the amount of sulfur consumed in the formation of metal sulfides. Based on the weight

of substrates (1.10 g), amount of sulfur in the jet fuel (1160 ppm), the fuel flow rate (1 mL/min)

and the elemental composition of iron and nickel for each substrate from Table 2.1, the atomic

percentages of iron, nickel and sulfur are calculated and summarized in Table 2.2. These values

are used to predict the sulfide phases observed in FESEM using the ternary phase diagram. The

dotted lines in the phase diagram corresponding to 35 at% sulfur and 55 at% Iron are shown for

convenience. They do not have any physical significance. Figure 2.4 shows the Fe-Ni-S phase

diagram used for this purpose. Figure 2.5 and 2.6 show the FESEM and TPO of deposits formed

on AISI 316 and AISI 321 respectively. As expected, Figure 2.5a shows the presence of metal

sulfides in the form of faceted crystallites denoted as (P) and fibers denoted as (F) on AISI 316.

The diameter of the metal sulfide fibers measured in the FESEM is found to be 75 nm. Similar

structures were observed during thermal stressing with a different batch of high-sulfur containing

jet fuel that was exposed to AISI 316 at 470 ºC [20]

. The phase diagram in Figure 2.4 predicts the

formation of pyrrhotites. FESEM shows the presence of faceted prismatic structures and

filamentous structures. The presence of this phase in two different morphologies was also

observed under pyrolytic conditions [20]

. The formation of pyrrhotites in the temperature range

250 ºC – 500 ºC on steel surfaces during the processing of crudes is observed in petroleum

refining [21]

. Therefore, it is suggested that these structures belong to the pyrrhotite phase.

42

The amount of carbon deposits on AISI 316 is marginally less than that on AISI 321. The

micrographs Figure 2.5a and 2.6a corresponding to AISI 316 and AISI 321 respectively show

that the extent of degradation in AISI 316 is less than that on AISI 321 due to the formation of

metal sulfides which increases the surface roughness during the course of experiment. As

mentioned above the iron content of AISI 316 is lower than those of AISI 304 and AISI 321.

This may be responsible for the latter showing more extensive metal sulfide formation and

roughening of substrate surfaces that create more area for carbon deposition. The carbonaceous

deposits seen as bright regions (A1) in Figure 2.5a are scattered along the length of the

filamentous pyrrhotite structures. The TPO profile for AISI 316 in Figure 2.5b shows three

groups of peaks. The broad spectrum in the low temperature range between 250 ºC and 400 ºC

can be attributed to relatively more reactive thermal deposits due to the formation of hydrogen-

rich carbonaceous solid from higher alkanes in Jet-A [7]

. These deposits are formed more likely

by liquid phase polymerization reactions and condensation reactions. The micrograph (Fig 2.5a)

also shows the presence of spherulitic carbon (S). These spherulitic carbon structures may have

formed by nucleation and growth of precursors in the gas phase followed by their deposition on

the surface of sulfides. High resolution transmission electron microscopy (HRTEM) examination

has shown that the spherulitic deposits accumulating on the sulfide particles are amorphous in

nature. The intermediate broad peak in the temperature range between 400 °C and 500 °C can be

attributed to the oxidation of spherulitic deposits [10]

. The formation of layers of carbonaceous

thin films and spherulitic deposits on the surface of metal sulfides has also been observed during

pyrolytic degradation of Jet-A at 470 ºC [10]

. The most intense high temperature peak seen in the

profile between 500 ºC and 700 ºC may be attributed to the carbonaceous film or platelets

43

formed on metal surfaces through dehydrogenative catalysis, producing greater structural order

in the solid carbon deposit in Figure 2.5b.

As seen in Figure 2.6b, the TPO profile for AISI 321 has three peaks that are better resolved

and have a more uniform distribution of peak intensities compared to the TPO profile of the

deposits on AISI 316. Similar to the suggested assignments for AISI 316 deposits, the first peak

(250 ºC and 400 ºC) can be attributed to the hydrogen rich-carbonaceous solid from

decomposition of higher alkanes. The intermediate peak (400 ºC – 500 ºC) may be attributed to

the oxidation of spherulitic carbon marked as S in Figure 2.6a and particulate deposits observed

as bright white regions on the surface of prismatic metal sulfide crystallites. The high

temperature peak between 500 and 700 ºC can be attributed to the oxidation of structurally more

ordered deposits. The phase diagram in Figure 2.4 predicts the presence of pyrrhotites. Surface

morphology of deposits observed in FESEM shows the presence of prismatic metal sulfides (P)

and filaments (F) in Figure 2.6a. Pyrrhotites are known to have these two morphologies.

Therefore, it is suggested that the crystallites on the surface of the AISI 321 are pyrrhotites.

The FESEM micrograph of Figure 2.7a shows the presence of filaments (F), faceted metal

sulfides (P) and spherulitic deposits (S) on AISI 304. X-ray diffraction of the sample containing

these deposits in Figure 2.7b shows the presence of hexagonal pyrrhotites having six- fold

symmetry that are observed in the micrograph marked as H1. The ternary phase diagram in

Figure 2.4 predicts the same. It is noteworthy to say that the signal was strong only from this

substrate. XRD on other substrates did not produce a good signal to detect the presence of

sulfides. The TPO profile in Figure 2.7c appears to contain two broad peaks. But upon closer

inspection, one may see that the profile can be deconvoluted to four peaks in the temperature

ranges 250 – 400 ºC, 400 – 600 ºC, 550 – 650 ºC and 550 – 720 ºC. These peaks may be assigned

44

in the order of decreasing reactivity, to hydrogen-rich carbonaceous deposits, spherulitic solid

carbon deposits, small particles of ordered carbons formed by metal catalysis, and large platelets

or films of ordered carbon structures, respectively.

Figure 2.8a shows the FESEM micrograph of deposits formed on AISI 347. The phase

diagram shown in Figure 2.4 predicted the presence of pyrrhotites which can be observed in the

form of prismatic structure marked as P and filamentous structure marked as F in the

micrograph. The amount of carbonaceous deposits on AISI 347 is 1.4 times greater than that on

AISI 321. The three broad regions of peaks observed on the TPO profile of AISI 347 in Figure

2.8b can be similarly assigned to hydrogen-rich carbonaceous deposits (200 – 450 ºC),

spherulitic solid carbon deposits (450 – 600 ºC), and structurally ordered deposits formed by

metal catalysis respectively. Compared to the TPO of AISI 304 and AISI 321 deposits, the

multiple high temperature peak(s) have shifted to higher temperatures, suggesting a stronger

catalytic activity of AISI 347 than other stainless steels, and is also evident in the larger amount

of carbon deposits formed on AISI 347. Given the only major difference in composition among

AISI 347, AISI 304 and AISI 321 is the presence of minor component Nb (1.2%) in AISI 347, it

is not clear as to what causes the catalytic activity.

Among all the iron-rich alloy surfaces selected for thermal stressing at 350 °C, FecrAlloy

gave the highest amount of carbon deposits. It is well known that iron and iron oxides catalyze

dehydrogenation reaction and carbon deposition [22]

. The above data suggest that the possibility

of sulfide formation should be considered. In the FESEM micrograph Figure 2.9a, it is

interesting to note the absence of metal sulfides and the presence of deposits formed from gas

phase denoted as E1, spherulitic deposits (S) and bright regions resembling structurally less

ordered carbon aggregates on the surface of the deposits (B). As the metal is an aluminized iron

45

and chromium alloy, passivated by chromium and aluminum, the sulfide formation is not

observed at 350 ºC. The TPO profile, Figure 2.9b shows the presence of three peaks having

similar intensities, resembling the case with AISI 321 (Figure 2.6b). The presence of aluminum

in the metal prevents severe degradation of the surface relative to other metals due to their

exposure to sulfur in jet fuel.

The behavior of Inconel 600 towards deposit formation is different from other alloys. The

amount of metal sulfides formed on Inconel 600 shown in the FESEM micrograph Figure 2.10a

is much less compared to iron-rich alloy surfaces that form metal sulfides. In spite of the lower

amount of metal sulfide formation, the metal shows a relatively higher catalytic activity towards

carbon deposition. The composition of the alloy (Table 2.1) shows the presence of copper as a

minor element. Studies conducted to elucidate the effect of metals on deposit formation from jet

fuel showed that among the metals copper, nickel, iron and cobalt, copper showed the highest

catalytic activity for deposit formation in the temperature range 200 °C – 260 °C [23]

. Therefore,

the presence of copper increases the catalytic activity of metal towards formation and/or

decomposition of oxidation products and hence carbon deposition. The oxidation of deposits

with different morphologies as observed in FESEM (Figure 2.10a) gives rise to CO2 peaks in the

temperature range between 370 °C and 600 °C in the TPO profile shown in Figure 2.10b. The

higher intensity of the high temperature peak and greater amount of carbon deposits can be

attributed to high catalytic activity of copper. The phase diagram predicts the formation of

heazlewoodite during the interaction of metals with sulfur compounds in jet fuel. The formation

of filamentous structures was observed on Inconel 600 under pyrolytic conditions [20]

. TEM

investigations of these filamentous structures suggest the presence of heazlewoodite [20]

.

Therefore, it is suggested that the fiber-like structures (F) obtained in this experiment and seen in

46

the micrograph (Figure 2.10a), are heazlewoodite crystals. The facts that the formation of metal

sulfides is less on Inconel 600 and the PFPD data suggest that sulfides and disulfides might

dissociate due to weak bonding between sulfur atoms and alkyl carbon atoms and accelerate the

chain initiation reactions that contribute to the formation of carbonaceous solids in the

intermediate regime.

The deposits on the surface of Inconel 718 shown in Figure 2.11a are structurally disordered

carbon seen as bright regions (B), spherulitic deposits (S) and sulfides in the form of fibers (F).

The sulfide crystals are suggested to be heazlewoodite crystals based on the phase diagram and

the above analysis for Inconel 600. Contrary to the observation under pyrolytic conditions, the

amount of solid carbonaceous deposits on Inconel 718 as shown in Figure 2.11b is greater than

that on Inconel 600. The presence of niobium, titanium and molybdenum in Inconel 718

appeared to suppress carbon deposition under pyrolytic conditions [7]

. In the intermediate

regime, where both pyrolysis and liquid phase autoxidation of hydrocarbons contribute to carbon

deposition, transition metals exhibit varying degree of catalytic activity toward the hydrocarbons

during dehydrogenation and carbon-carbon bond cleavage in catalytic cracking. The

participation of metals with multiple valence states has also been observed in the liquid phase

autoxidation of hydrocarbons. The metals catalyze the reduction-oxidation reaction of

hydroperoxides, by forming metal-hydroperoxide complex [24]

, and other products formed during

the oxidation of hydrocarbons, resulting in the formation of free radicals [25]

. Iron, cobalt, nickel,

copper, chromium and manganese are generally used as catalysts during oxidation of

hydrocarbons to accelerate the reduction-oxidation reaction [25]

. Sulfur compounds under

pyrolytic conditions are known to either promote or inhibit carbon deposition depending upon

the type of metal substrate [26]

. Under pyrolytic conditions, addition of thiophenes, 3-methyl

47

benzothiophenes to n-dodecane inhibited carbon deposition on iron and nickel surfaces by

blocking the active sites, whereas benzyl phenyl sulfide promoted the formation of carbonaceous

solids on Inconel 718 [26]

. The above facts along with the PFPD data help to speculate the reason

for the formation of more deposits on Inconel 718 relative to Inconel 600.

From Figure 2.3, in the intermediate regime under consideration, it can be seen that among

the Inconel alloys used, Inconel 750-X gives the highest amount of deposits. The FESEM

micrograph Figure 2.12a shows the presence of disordered carbon seen as bright regions (B) on

the surface of prismatic metal sulfides and spherulitic deposits (S). The TPO profile Figure 2.12b

shows that the structurally less ordered carbon oxidizes to give a broad plateau in the

temperature range between 300 °C and 480 °C and the relatively ordered deposits oxidize at

approximately 600 °C. The formation of metal sulfides in Inconel 750X was not observed under

pyrolytic conditions with JP-8 when the concentration of sulfur was 68 ppm [27]

. It was also

observed that metal surfaces exposed to Jet-A had significant degradation at 500 °C as opposed

to JP-8. As the concentration of sulfur is 1160 ppm in the Jet-A sample used in these

experiments, metal sulfides in the form of fibers (F) and prismatic crystallites (P) formed on the

substrate are observed in the micrograph. The phase diagram predicts the formation of

pentlandite and heazlewoodite. Filamentous structures observed in the micrograph are suggested

to be heazlewoodite crystallites based on the analysis for Inconel 600. As inconel 750-X is a

nickel rich alloy, heazlewoodite would have formed during the early stages of the reaction.

Subsequently monosulfide solid solution might have formed due to the reaction between

organosulfur compounds and the minor element iron. Pentlandites typically form by the reaction

between monosulfide species and heazlewoodite.

48

To verify the catalytic activity of metals at 350 ºC, thermal stressing was conducted with

Silicon substrate. The morphology and amount of the deposits formed on Silicon are shown in

Figure 2.13a and 2.13b respectively. The structurally disordered carbonaceous deposits seen as

bright white regions (B) and the spherulitic deposits (S) shown in Figure 2.13a oxidize in the

temperature range between 250 ºC and 500 ºC as shown in the TPO profile in Figure 2.13b. The

absence of high temperature peaks at temperatures greater than 500 ºC suggests that iron and

nickel rich alloys exhibit catalytic activity for the formation of carbonaceous deposits during

thermal stressing at 350 ºC. The absence of metal sulfides on the surface of Silicon and the lower

amount of carbon deposits also suggest that the formation of metal sulfides on iron and nickel

rich alloys increase the surface area available for carbon deposition.

2.5 Conclusions

The formation of carbonaceous solid deposits on metal substrates in the intermediate

regime is influenced by reactive organic sulfides and disulfides in the jet fuel, decomposition of

oxidation products from liquid phase autoxidation, pyrolysis along with metal catalysis and the

metal sulfide formation. Characterization of the carbonaceous deposits by FESEM and TPO

shows predominantly the presence of spherulitic deposits which nucleate and grow in the fluid

phase. The formation of metal sulfides increases the surface roughness and causes disruption of

the surface significantly. Based on the characterization and the prediction from phase diagram, it

appears that pyrrhotite forms in the iron rich metals and heazlewoodite forms in the nickel rich

alloy surfaces. Due to the surface disruption, metals with multiple valence states may be exposed

to the oxygenated intermediates and participate in the decomposition of hydroperoxides through

metal-hydroperoxide complex and other oxidation products formed during liquid phase

autoxidation. Therefore, the amount of carbon deposition on the alloys increased in the

49

following order AISI316 < AISI 321 ~ AISI 304 < Inconel 600 < AISI 347 < Inconel 718 <

FecrAlloy < Inconel 750-X. The presence of molybdenum, titanium and niobium in smaller

amounts does not appear to affect carbon deposition under the experimental conditions. Carbon

deposition on FecrAlloy, Inconel 600, Inconel 718 and Inconel 750-X shows that the formation

of metal sulfides do not necessarily passivate the surface and reduce carbon deposition.

50

2.6 References

1. Taylor, W. F., Kinetics of Deposit Formation from Hydrocarbons - Fuel Composition Studies.

Industrial & Engineering Chemistry Product Research and Development 1969, 8 (4), 375-380.



40 (2), 596-603.

3. Altin, O.; Eser, S., Analysis of Carboneceous Deposits from Thermal Stressing of a JP-8 Fuel

on Superalloy Foils in a Flow Reactor. Industrial & Engineering Chemistry Research 2001, 40

(2), 589-595.

4. Jones, E. G.; Balster, W. J.; Balster, L. M., Aviation Fuel Recirculation and Surface Fouling.

Energy Fuels 1997, 11 (6), 1303-1308.

5. Jones, E. G.; Balster, L. M.; Balster, W. J., Thermal Stability of Jet-A Fuel Blends. Energy

Fuels 1996, 10 (2), 509-515.

6. Taylor, W. F., Kinetics of Deposit Formation from Hydrocarbons .3. Heterogeneous and

Homogeneous Metal Effects. Journal of Applied Chemistry of the Ussr 1968, 18 (8), 251-&.

7. Venkataraman, R.; Eser, S., Characterisation of solid deposits from the thermal-oxidative

degradation of jet fuel. Int. J. Oil, Gas and Coal Technology 2008, Vol. 1, Nos. 1/2, 126-137.



Research 2006, 45 (26), 8946-8955.

9. Zhang, F. Carbon Deposition on Heated Alloy Surfaces from Thermal Decomposition of Jet

Fuel. M.S. Thesis, The Pennsylvania State University, University Park, PA 16801, 2000.

10. Song, C.; Ma, X. L., New design approaches to ultra-clean diesel fuels by deep

desulfurization and deep dearomatization. Appl. Catal. B-Environ. 2003, 41 (1-2), 207-238.



2008, 47 (23), 9337-9350.

12. Hazlett, R. N.; Hall, J. M.; Matson, M., Reactions of Aerated N-Dodecane Liquid Flowing

over Heated Metal Tubes. Industrial & Engineering Chemistry Product Research and

Development 1977, 16 (2), 171-177.



fuel. Energy Fuels 2003, 17 (3), 577-586.

51

14. Venkataraman, R. Solid Deposit Formation from the Pyrolytic and Oxidative Degradation of

Jet Fuel and Diesel Fuel Ph. D Thesis, The Pennsylvania State University, University Park, PA

16801, 2007.

15. Querini, C. A.; Fung, S. C., Temperature-Programmed Oxidation Technique - Kinetics of

Coke O-2 Reaction on Supported Metal-Catalysts. Appl. Catal. A-Gen. 1994, 117 (1), 53-74.

16. Eser, S.; Venkataraman, R.; Altin, O., Utility of Temperature-Programmed Oxidation for

Characterization of Carbonaceous Deposits from Heated Jet Fuel. Industrial & Engineering

Chemistry Research 2006, 45 (26), 8956-8962.

17. Link, D. D.; Baltrus, J. P.; Rothenberger, K. S., Class- and structure-specific separation,

analysis, and identification techniques for the characterization of the sulfur components of JP-8

aviation fuel. Energy Fuels 2003, 17 (5), 1292-1302.

18. Trimm, D. L., Fundamental Aspects of the Formation and Gasification of Coke. In Pyrolysis-

Theory and Industrial Practice, Albright, L. F., Crynes, B. L., Corcoran, W. H. Academic Press:

1983; p 229.

19. Raghavan, V., Fe-Ni-S (Iron-Nickel-Sulfur). J. Phase Equilib. Diffus. 2004, 25 (4), 373-381.



Research 2008, 47 (23), 9351-9360.

21. Pareek, V. K.; Ozekcin, A.; Mumford, J. D.; Ramanarayanan, T. A., Transport of sulfur

through preformed spinel films on low alloy Fe-Cr steels. J. Mater. Sci. Lett. 1997, 16 (2), 128-

130.

22. Tanabe, K.; Misono, M.; Ono, Y. n; Hattori, H., New Solid Acids and Bases. Their Catalytic

Properties. Studies in surface science and catalysis. Elsevier: 1989; Vol. 51, p 60-64.

23. Taylor, W. F.; Wallace, T. J., Kinetics of Deposit Formation from Hydrocarbons - Effect of

Trace Sulfur Compounds. Industrial & Engineering Chemistry Product Research and

Development 1968, 7 (3), 198-&.

24. Hazlett, R. N., Thermal Oxidation Stability of Aviation Turbine Fuels. ASTM 31 -001092-

12: Philadelphia, 1991.

25. Emanuel, N. M.; Denisov, E. T.; Maizus, Z. K., Liquid-Phase Oxidation of Hydrocarbons.

Academy of Sciences of the USSR 1967.

26. Raymundo-Pinero, E.; Altin, O.; Eser, S., Effect of Sulfur Compounds on Solid Deposition

on Metals and Inconel 718 from Thermal Decomposition of N-Dodecane. Preprints-American

Chemical Society, Division of Petroleum Chemistry 2001, 47 (3), 216-218.

52

27. Altin, O.; Eser, S., Characterization of carbon deposits from jet fuel on Inconel 600 and

Inconel X surfaces. Industrial & Engineering Chemistry Research 2000, 39 (3), 642-645.

53

Table 2.1. Elemental Composition of Alloys. (Goodfellow Ltd)

Substrate Elemental Composition

Fe Ni Cr Mn C Ti Mo Si Nb S Cu Al Y Zr

AISI316

AISI321

AISI304

AISI347

FeCrAl

IN600

IN718

IN750-X

64

68

69

67

72.6

8

18.5

7

12

10.5

10

11

74.43

52.5

73

18

18

18

18

22

15.5

19

15.5

2

2

2

2

1

0.18

0.5

0.08

0.08

0.08

0.08

0.03

0.15

0.04

0.04

0.6

0.9

2.5

3

3

1.0

0.5

1.0

1.0

0.3

1.0

0.18

0.25

1.2

5.13

0.95

0.03

0.03

0.03

0.03

0.0015

0.0008

0.0005

0.5

0.15

0.25

4.8

0.5

0.7

0.1

0.1

Table 2.2. Calculated Atomic Percentage of Fe, Ni, and S on the Alloys after 5 hours of

Thermal Stressing.

Substrate Elements considered for sulfide formation

Fe Ni S

AISI316

AISI321

AISI304

AISI347

IN600

IN718

IN750-X

54

57

57

56

7

17

6

10

8

8

9

59

45

58

36

35

35

35

35

38

36

54

Figure 2.1. Flow reactor setup for thermal stressing experiment.

a)

55

b)

Figure 2.2. (a) GC-MS chromatogram of Jet-A showing the composition of the fuel (b) PFPD

chromatogram showing sulfur compounds.

Figure 2.3. Carbon deposits on different metal surfaces from Jet-A at 350 °C, 500 psig with a

fuel flow rate of 1 mL/min for 5 h.

56

Figure 2.4. Fe-Ni-S ternary phase diagram at 400 ºC - adapted from (Raghavan, 2004).

Terminology – γ-Continuous solid solution between face centered cubic iron and nickel, py-

pyrite FeS2, pn-pentlandite (FeNi)9S8, hz-heazlewoodite Ni3S2, mss-monosulfide solid solution,

vio-violarite Ni3S4, vs-vaesite NiS2, α-iron rich region

.


A at 350 °C and 500 psig for 5 h.

b a

P

F

S

A1

57



Figure 2.7. (a) FESEM image (b) X-ray diffractogram and (c) TPO profile of the deposits

formed on AISI 304 from Jet-A at 350 °C and 500 psig for 5 h.

a

a

b

c

b

b

c

b

F

P

S

F

H1

S

F

P

58



Figure 2.9. (a) FESEM image and (b) TPO profile of the deposits formed on FecrAlloy from Jet-


a

b

b

E2

E3

E1

a

F

P

a S

E2

S

B

E1

a

59


Jet-A at 350 °C and 500 psig for 5 h.


Jet-A at 350 °C and 500 psig for 5 h.

a

a

b

b

F

B F

S

S

60

Figure 2.12. (a) FESEM image and (b) TPO profile of the deposits formed on Inconel 750X

from Jet-A at 350 °C and 500 psig for 5 h.

Figure 2.13. (a) FESEM image (b) TPO profile of the deposits formed on Silicon from Jet-A at

350 °C and 500 psig for 5 h.

a b

S

F

P

B

B

S

a b

61

Chapter 3

Environmental Barrier Coatings by MOCVD on tube surfaces to inhibit carbon deposition

3.1 Background

Environmental barrier coatings play an important role by extending the life of components

exposed to high temperatures and harsh environments and improving the reliability and

performance of these components. Thermal barrier coatings were widely used in turbines for

propulsion and power generation to insulate the underlying structure against the effect of heat,

retard creep degradation and reduce the severity of thermal transients. Higher operating

temperatures increase the efficiency of turbines and serve as the driver for the development of

materials. Cooling of turbine blades coupled with TBCs and improvements in metallurgical

composition of the alloys had increased the operating temperatures to 90% of the melting point

of alloys. In spite of these developments, the application of superalloys at high temperatures is

limited by their melting point which has led to the search for alternatives. The very high melting

point, high temperature strength and low density make ceramic structures like SiC and Si3N4 a

potential candidate for turbine materials used in power generation. But exposure of these ceramic

materials to water vapor present in the working fluid at high temperatures results in the

formation of volatile Si(OH)x species. Therefore recent development has focused on coatings

that can play a dual role where they can provide thermal insulation to the underlying metal and

be an environmental barrier to avoid any degradation caused by the interaction of the metal with

the working fluid [1]

. To enhance the ductility and toughness of ceramics, the development of

ceramic matrix nanocomposites system is underway. Intermetallics like nickel aluminide (NiAl)

and titanium aluminide (TiAl) are researched as matrix materials for composites. Low creep

strength at high temperatures and poor ductility at lower temperatures are the shortcomings in

62

intermetallics compounds. Alloying of NiAl to improve the mechanical properties might result in

the loss of the excellent oxidation-resistance of near stoichiometric NiAl. Therefore coatings

may, thus, be necessary to provide the required oxidation resistance for intermetallics as well.

As a general rule, the coatings should improve the system durability and reliability. The

operating temperature of a turbine and the cyclic capability are inversely related to one another

[3]. The interaction between high temperature creep of materials and thermomechanical fatigue

determines the life expectancy of environmental barrier coatings [2]

. A barrier coating can reduce

the magnitude of thermomechanical fatigue. The selection of a barrier coating depends upon the

dedicated purpose of the turbine, size and rotational speed, number of on/off duty cycles, time-

at-temperature and the quality of the fuel used. Barrier coatings for power turbines are subjected

to longer operational hours between overhauls, higher operational temperatures, fewer on/off

duty cycles and aggressive ambient characterized by the presence of ammonia (NH3), hydrogen

sulfide (H2S), hydrogen chloride (HCl) and few parts per million of trace alkali metals

environment. Even though these corrosive chemicals are removed significantly before the fuel

enters the turbine, the coatings protect various metal components from trace amounts of H2S and

NH3 in the environment, and reduce the driving force for creep and fatigue [3]

. Although, the

coatings do not reduce the transport of oxygen to the substrate, the temperature of the substrate is

lowered by various cooling configurations of the blades and hence the oxidation rate due to hot

corrosion is also reduced [2]

. The operating conditions of aircraft turbines are characterized by

frequent on/off duty cycles and lower operational temperatures. The barrier properties of the

coatings depend upon the location where the coatings are intended to be applied. The fuel

handling system of the aircraft engine handles sulfur-containing jet fuel and experiences lower

temperatures. The turbine blade handles combustion products containing sulfur dioxide at high

63

temperatures. In this chapter, the material and process development for barrier coatings will be

discussed to select a process that can be used for the investigation of environmental barrier

coatings in the fuel handling system of aircraft engines. The selected process is further evaluated

by coating tubes of varying diameters with different coating precursors at different conditions

and the coatings are characterized.

3.2. Coating process for EBCs:

The various atomistic and particulate deposition methods and their benefits and limitations are

summarized in Table 3.1. The three methods used predominantly for the deposition of barrier

coatings on the surface of turbine blades are plasma spray deposition, electron beam physical

vapor deposition and electrodeposition.

3.2.1. Plasma Spray Deposition

Atmospheric plasma spray coating is a common method to form thermal barrier coatings

on the surface of turbine blades. The typical configuration of a plasma spray coating system is

shown in Figure 3.1. Aluminum oxide and zirconia partially stabilized with yttria are widely

deposited by this method for industrial applications. The material to be coated is fed in the form

of a powder at a constant rate. Plasma is generated by passing a mixture of inert argon gas and an

enthalpy enhancing gas like hydrogen between the copper anode and the tungsten cathode in a

plasma torch. The electrodes in the plasma torch are water cooled. A very high voltage applied

between the electrodes produces an electric arc that ionizes the gas which exits the nozzle. The

recombination of ions releases the enthalpy yielding a temperature as high as 15000 K for a

plasma torch with a power of 40 kW enough to melt any ceramic material. The powders are fed

behind the nozzle with a carrier gas. They are heated, melted and partially evaporated in the

64

plasma jet. The substrate is placed in the line of sight of the plasma jet separated by a certain

distance to deposit ceramic films. The plasma torch moves in front of the substrate at a certain

velocity of few mm per second. The powder particles undergo rapid cooling at a rate of 104

K/sec and solidification. Typically, plasma spray coatings are more than 50 μm in thickness.

Figure 3.1. Atmospheric plasma spray method for TBCs [4]

In order to approach the theoretical density and extremely high adhesion strength, metal bond

coats are deposited by low pressure plasma spray (LPPS). It is competitive with electron beam

physical vapor deposition (EBPVD) because of the compositional flexibility and high deposition

rates. LPPS is restricted to line-of-sight deposition process. The substrate in plasma spray

coating is heated to 900 °C – 1000 °C. The turbine blade to substrate distance is about 10 in. – 16

in. in the chamber and the pressure is 30 – 60 Torr. The power of the gun is 80 kW. The powder

feed rate can vary from 3 to 20 kg/hr depending upon the application.

Particles that melt in the plasma take the shape of a spherical droplet to minimize the

surface energy and upon impingement on the substrate deposits layered structures [4-6]

. The

structure of each layer containing molten and solid phase results is metastable. When annealed at

the crystallization temperature, the porosity decreases and the layered structure changes to

globular structure. The density of plasma sprayed coating falls between 85% and 93% of the

65

theoretical density of the same but more compact materials [4]

. The evolution of gases like

hydrogen, nitrogen and oxygen during cooling due to the reduced solubility results in open

porosity and closed porosity. The enclosed voids cannot be removed by changing the coating

conditions. Sprayed coatings are more brittle than the corresponding compact materials. The

adherence of all sprayed coatings to the substrate reduces with thickness. Addition of subsequent

layers adds internal stress to the sprayed coating. When the internal stress exceeds the bond

strength of the coating, delamination of the coating occurs. Cooling the substrate during plasma

spray coating is recommended to reduce the internal stress of the coating.

The plasma spray coating has also been used to coat the interior surfaces in the aircraft

industry by mounting the gun in a pole to extend it physically into long tubes. It has been

observed that the deposition efficiency of this method is 65% for metal oxides and 80% for

metals [7]

. In other words, 35% of the metal oxide and 20% of the metals termed as overspray are

not coated as they remain as dust in the interior of the conduit without being incorporated in the

coating. The overspray coats the spray gun, dampening the arc from the plasma spray gun and

subsequently interrupts the coating process by preventing arc initiation. The confinement of the

overspray inside a conduit heats them causing undesired physical adhesion over the coating.

Even though new devices are built to coat the internal surfaces that can reduce the overspray by

providing a volume of reduced pressure to draw the overspray from the interior of the conduit [8]

,

these coating methods are not suitable for tubes with diameter as small as 1.5 mm ( 1/16 of an

inch) that are typically used in the fuel injectors.

66

3.2.2 Electron Beam Physical Vapor Deposition (EB-PVD)

In an EBPVD system, the deposition chamber is evacuated to a pressure of 10-4

Torr. The

materials to be evaporated are in the form of ingots. There are as many as six electron guns, each

having a power from few tens to hundreds of kW. Electron beam generated by thermionic

emission, field emission or the anodic arc method is accelerated to a high kinetic energy and

focused towards the ingot. When the accelerating voltage is between 20 kV – 25 kV and the

beam current is a few amperes, although some of incident electron energy is lost in the excitation

of X-rays and secondary emission, 85% of the kinetic energy of the electrons is converted into

thermal energy as the beam bombards the surface of the ingot. The surface temperature of the

ingot increases resulting in the formation of a liquid melt. The liquid ingot material evaporates

under vacuum. The ingot itself is enclosed in a copper crucible cooled by water circulation. The

level of molten liquid pool on the surface of the ingot is kept constant by vertical displacement of

the ingot. The number of ingot feeders depends upon the material to be deposited. The

evaporation rate may be of the order of 10-2

g/cm2 sec

[9].

Direct evaporation, reactive evaporation and co-evaporation are the three methods used to

deposit coatings for industrial applications [9]

. Carbides like titanium carbide and borides like

titanium boride and zirconium boride that can evaporate without decomposition in the vapor

phase are compacted in the form of an ingot, evaporated in vacuum by the focused high energy

electron beam and the vapors are directly condensed over the substrate.

Certain refractory oxides like Al2O3, carbides like SiC and WC undergo fragmentation

during their evaporation by the electron beam and form non-stoichiometric species [9]

. These

compounds can be deposited on the substrate either by reactive evaporation or by co-

http://en.wikipedia.org/wiki/Carbide

http://en.wikipedia.org/wiki/Boride

67

evaporation. In the reactive evaporation process, the metal is evaporated from the ingot by the

electron beam. The vapors are carried by the reactive gas, which is oxygen in case of metal

oxides or acetylene in case of metal carbides. When the thermodynamic conditions are met, the

vapors react with the gas in the vicinity of the substrate to form films. Metal carbide films can

also be deposited by co-evaporation. In this process, two ingots are used, one for metal and the

other for carbon. Each ingot is heated with different beam energy so that their evaporation rate

can be controlled. As the vapors arrive at the surface, they chemically combine under proper

thermodynamic conditions to form a metal carbide film [9]

.

EBPVD systems are equipped with ion sources. These ion sources are used for substrate

etching and removal of coatings on unwanted areas, sputtering the target and controlling the

microstructure of the substrate. The ion beams bombard the surface and alter the microstructure

of the film. When the deposition reaction takes place on the hot substrate surface, the films

develop an internal tensile stress due to the mismatch in the coefficient of thermal expansion

between the substrate and the film. The residual stress in the films can be minimized by the

proper choice of coating conditions, coating materials and substrate [10]

. High energy ions can be

used to bombard these ceramic thermal barrier coatings and change the tensile stress into

compressive stress. Ion bombardment also increases the density of the film, changes the grain

size and modifies amorphous films to polycrystalline films.

Under the high vacuum conditions existing in EBPVD, the mean free path of the molecules

between successive collisions is much greater than the source-to-substrate distance. The

trajectory of the vapor molecules is set by the evaporation of ingots by the electron beam. This

process has been used predominantly to coat the outer surfaces of turbine airfoils by the

translational and rotational motion of the complex geometry in the generated vapor cloud [9]

. The

68

outer surface that is not directly above the line-of-sight of the source receives less vapor flux. As

the incident angles of the vapor flux are inclined to the substrate in these non-line-of-sight

regions, columnar structure is not obtained. This disadvantage is surmounted by the use of

directed vapor deposition [11]

. In this process, the flow of carrier gas at supersonic conditions

carries the vapor cloud to the regions that occupy the non-line-of-sight locations [12]

. The

following idea is proposed to modify the EBPVD system to coat the interior surface of a

complex geometry. One end of the geometry should be left open to the system so that the carrier

gas with the vapor cloud can pass through the internal surface of the geometry that has to be

coated. The other end of the geometry can be connected to a vacuum pump with flexible metal

hoses. The geometry can be attached to the translating arm of the EBPVD system so that the

open end of the geometry can be placed near the vapor plume emanating from the source. The

pressure gradient created by the vacuum pump helps the carrier gas with the vapor cloud to flow

through the interior surface of the complex geometry before leaving the system. This approach

has not been used in any EBPVD system till date. Even though, an EBPVD system can be used

in coating complex geometries with these modifications, this process is not pursued further due

to the poor material utilization efficiency of 3% and high cost involved in the installation and

operation of the system.

3.2.3 Electrodeposition Process

The electrodeposition process used for the deposition of functional ceramic coatings can

be classified into two types based on the raw materials used. Electrophoretic deposition process

uses a suspension of ceramic particles for the deposition of a thick film. The particles are

suspended in a non-aqueous solvent typically ethyl alcohol to avoid gas formation. Phosphate

ester is typically used to disperse the particles by electrostatic stabilization. These dispersants

69

donate a proton to the surface of the ceramic particles and give them a positive charge so that

they can migrate in the presence of electric field in cathodic electrophoretic deposition. Polyvinyl

butyral is used as a binder in the suspension to increase the strength and adherence of the

deposited material and prevent cracking [12]

. The stoichiometry of the deposit depends upon the

stoichiometry of the powder used in the suspension. This process can be conducted in constant-

current and constant-voltage regime. The deposition rate depends upon the magnitude of applied

electric field, concentration of the suspension and the electrophoretic mobility of particles. The

uniformity and adhesion of deposits can be improved by the use of electrolytes. But there is a

critical electrolyte concentration below which the suspension is stable and above which

aggregation and sedimentation of particles take place. Particle aggregation results in the

formation of a porous coating. To reduce the deposition time, higher voltages are applied

between electrodes. The deposition rate of this process can be between 1 and 103 mm per minute.

This process was used for the deposition of β-alumina on the outer surface of a graphite tube for

the application of sodium-sulfur batteries from a suspension of alumina particles in

dichloromethane and trichloroacetic acid [13]

. The alumina particles in this suspension are

negatively charged. By using the tube or the coating geometry as a positive electrode, the slurry

can be pumped into the geometry so that thick uniform deposits can be obtained on the inner

wall of stainless steel tube within short duration of time. The coating obtained will be porous and

sintering of the coating at high temperatures is necessary to eliminate pores between particles in

the coating. Exposing the geometry to be coated to such high temperatures is undesirable as it

can change the mechanical properties of the metal substrate. Therefore this coating method is not

pursued further.

70

3.2.4 Chemical Vapor Deposition (CVD)

Chemical vapor deposition (CVD) has been used for the development of bond coats for

TBCs [2]

, dielectric oxides and metallization interconnects used in semiconductor fabrication [14]

.

It can be used in the deposition of borides, carbides, nitrides, silicides and oxides for hard

coatings, protection against corrosion and diffusion barriers [15]

. One of the advantages of

chemical vapor deposition is its applicability to complex geometries that are in need of a non-

line-of-sight deposition process. The following key steps are involved in a CVD process [15]

.

1. Transport of precursor reactants into the gaseous reaction chamber.

2. Diffusion of the precursor from the gas phase across the boundary layer.

3. Adsorption of reactants and surface reactions followed by surface diffusion that controls

the microstructure of the films on the substrate.

4. Diffusion of gaseous products through the boundary layer, transport and removal of

gaseous products from the deposition chamber.

A fuel injector is an example of a complex geometry used in the fuel handling system of an

aircraft engine that consists of conduits of varying diameters from the inlet to the outlet. In a

conventional chemical vapor deposition process, diborane or boron trichloride, methane,

ammonia and silane react with metal halides to form the borides, carbides, nitrides and silicides

respectively. These deposition reactions take place at temperatures between 800 °C – 1200 °C

[15]. Exposing the geometry at such high temperatures to corrosive products like hydrogen halides

may cause degradation. Presence of a metal organic substance in the liquid form like aluminum

tri-secondary butoxide or in a solid form like aluminum acetylacetonate facilitates the precursor

decomposition at much lower temperatures and prevents the formation of corrosive products. In

71

the case of a solid or a liquid precursor, by sublimation or evaporation respectively, the precursor

is transported from the bubbler into the preheating line. The precursor is diluted in the gas phase

and transported into the reactor where the geometry to be coated is heated to the appropriate

reaction temperature. Therefore metalorganic chemical vapor deposition (MOCVD) is the

method chosen for the deposition of metal and ceramic coatings.

3.2.5 Effect of Process Variables on the MOCVD and Properties of Coatings

The primary property of an environmental barrier coating is to prevent the exposure of

the substrate to corrosive chemical environments and eliminate any undesired chemical reaction

between the working fluid and the substrate that would involve the consumption of the substrate

material leading to the degradation of its mechanical properties and improve the service life of

the substrate. To prevent the access of corrosive species to the substrate, the coating should be

free of pores. The coating has to be uniform and cover the entire geometry. In the combustion

chamber of a gas turbine used in power generation, apart from hot corrosion resistance, erosion

resistance and thermal insulation are the other properties necessary for a good environmental

barrier coating [15]

. As a fuel injector nozzle used in the aircraft gas turbine engine handles

aviation turbine fuel, erosion resistance for these coatings is not necessary. Thermal insulating

property of the coating used in the fuel injectors of an aircraft engine will be an additional

benefit derived from the thickness of the coating and appropriate material selection. Thermal

conductivity and thickness of the coating are the two characteristics that control their insulating

properties. Pressure, reactant to oxidant ratio, flow rate of reactant, flow rate of carrier gas,

substrate temperature and bubbler temperature are the principal variables that affect the growth

rate of a coating. Pressure affects the diffusion and mean free path of molecules in the gas phase.

Lower pressure enhances the diffusion of reactants across the boundary layer and improves the

72

conformal step coverage of high aspect ratio features. The ratio of reactant to carrier gas controls

the concentration of the reactant in gas phase. The composition of the carbide, oxide and nitride

coating can also be varied by the reactant to oxidant ratio. Excess reactant in the gas phase

causes all the reactions in the gas phase that would give rise to cauliflower morphology with a

large number of pores [16]

. The presence of pores is undesirable in protective surface coatings

because the sulfur compounds and the hydrocarbons in the fuel can access the metal substrate

that would lead to the formation of metal sulfides and the formation of carbonaceous deposits

due to the catalytic activity of metals. Peclet number and Damkohler number are the two

characteristic numbers that influence the distribution of thickness as a function of position of

sample in the reactor. Peclet number is defined as the ratio of convective flux to diffusive flux.

Damkohler number is defined as the ratio of rate at which a chemical species reacts at the surface

to the rate at which the reactive species diffuse onto the surface across the boundary layer.

Higher carrier gas flow rate increases the Peclet number to values greater than one indicating that

convective flux is greater than diffusive flux. Peclet number is independent of pressure. Lower

pressure increases the diffusive rate relative to the surface reaction rate and increases the

Damkohler number. Lower pressure and higher carrier gas flow rate are needed for the formation

of a uniform coating along the length of the tubes used in the fuel injectors.

Substrate temperature has a significant effect on the deposition profile of the coating along

the length of a geometry, composition and stress in the coating. When the carrier gas flows

through a conduit to coat its inner surface, the reactants have to diffuse to the wall of the reactor

that is perpendicular to the direction of convective flow of the carrier gas. Depending upon the

substrate temperature, the reactor can be operated in reaction rate-controlled regime or mass

transport-limited regime. Slower consumption of reactants in the reaction-rate controlled regime

73

at higher carrier gas flow rates may produce a thin coating along the length of the reactor at

lower substrate temperature. Faster consumption of reactants at higher substrate temperature in

the mass transport-limited regime increases the surface diffusivity and surface reaction rate to

produce a denser and thicker coating during a short duration of time in relatively a smaller

portion of the reactor. Therefore, the profile of the deposition rate can be flattened or narrowed

down from the inlet to the outlet of the geometry by controlling the flow rate and substrate

temperature. The presence of contaminants in the film introduced by metalorganic or

organometallic precursors changes the composition of the coating. Composition and density

affect the refractive index of the coating, optical and electrical properties of the film. Degassing

of impurities from the coating during each thermal cycle might affect the stability and

mechanical properties of environmental barrier coatings depending upon their operating

temperature.

3.2.6 Influence of Process Parameters in the Stress induced in Coatings

Residual stress can be described as a self-equilibrating internal stress that is introduced

during the process of deposition of a thin film on the substrate and remains in the system after

deposition. Intrinsic residual stress can be introduced in the film due to the presence of pores,

impurities left due to the incomplete decomposition of the precursor and partial growth of grains

due to high deposition rates or lower substrate temperature. Two mechanisms are proposed to

explain the nature of intrinsic stress in the CVD grown films [17]

. In adsorption-limited reaction,

gas-phase reactions results in a weaker bonding at the film-substrate interface and exerts an

upward force constrained by atomic packing and interface bonding and hence introduces intrinsic

compressive stress [17]

. In thermally activated reaction, agglomeration of islands by surface

diffusion for the lateral growth of a thin film introduces tensile intrinsic stress [17]

. Extrinsic

74

residual stress is caused due to the mismatch in the coefficient of thermal expansion between the

film and the substrate, the mismatch in the lattice parameter between the film and the substrate

and phase change of the film [18]

. The total residual stress is the sum of intrinsic and extrinsic

residual stress values. The residual stress in the coating does not generate any normal or shear

stress on flat, smooth and infinitely large surface. But the magnitude of these stress values

depends upon the geometry of the substrate in complex engineering components [19]

. Deposition

of a ceramic coating whose coefficient of thermal expansion is lower than that of the metallic

substrate results in the formation of compressive thermal stress in the film during cooling and

tensile stress during heating [20]

. For lower deposition temperature relative to the melting point of

the substrate, the intrinsic stress in the coating is higher than the thermal stress, whereas the

thermal stress exceeds the intrinsic stress values at higher deposition temperatures [21]

.

3.2.7 Coating Precursors

Chemical vapor deposition has been successfully used in the deposition of refractory

materials like titanium boride, silicon carbides, alumina and zirconia. Deposition of titanium

boride takes place at 1100 °C by the chemical reaction between titanium chloride, boron

trichloride and hydrogen [22]

. Silicon carbide is deposited by the decomposition of

monomethyltrichloro silane at 1000 °C in presence of pure hydrogen [23]

. Deposition of these

refractory materials on tubes of varying diameters and other complex geometries like a fuel

injector would expose them to these high temperatures and corrosive reaction products and is

therefore undesirable due to the degradation caused by the exposure of the substrate to the

corrosive reaction products at high temperatures. A metalorganic substance that does not contain

halogen is a preferable precursor so that the substrate is not exposed to halide vapors during

75

deposition. In order to conduct the process at low temperature and evaluate its feasibility to coat

complex geometries, alumina was chosen as a candidate.

3.2.8 Configuration of the MOCVD Experimental Set-up for coating tubes

A description of the components used in the MOCVD setup is shown in Table 3.2. A

schematic diagram of the MOCVD experimental set-up is shown in Figure 3.2. It consists of

three mass flow controllers numbered as 1, 2 and 3 shown in Figure 3.2 purchased from Omega

instruments and used for the discharge of carrier gas argon, purge argon, and ultra high purity

oxygen respectively. The bellow sealed valves used in the system are vacuum compatible and

perform the on-off function. A bubbler is custom-built with two manifolds. One end of the

manifold is connected to the outlet of the bellow sealed valve and the other end is connected to

the preheating line. The bubbler is typically washed with soap solution, acetone and ethanol and

then dried at 100 °C for 5 hours in an oven. It is then placed in a glove box which is flushed

with argon for five hours to eliminate any residual air or moisture inside the glove box. The

precursor is stored in a glove box in an argon environment and transferred into the bubbler inside

the glove box. The bubbler is then connected to the manifolds of the system. The bubbler is

heated to the desired temperature with a heating tape to keep the precursor in the vapor phase by

evaporation of the liquid precursor or sublimation of the solid precursor. The precursor vapors

are transported into the preheating line by the carrier gas argon. The bellow sealed valve marked

as number 5 in Figure 3.2 helps to eliminate the bubbler from the loop during cooling at the end

of the coating process. The residual vapors are removed with the help of purge argon which

flows through the mass flow controller 2. This line was also used to dilute the concentration of

the precursor vapor in the initial trails to optimize the process. Ultra high purity oxygen is used

as an oxidant in some cases and is fed to the preheating lines. The preheating line is maintained

76

at a temperature that is 50 °C higher than the bubbler temperature to avoid condensation of the

precursor vapors. The bubbler, preheating lines and exhaust lines are electrically heated by

heating tapes which are controlled by a seven segment CN1507 temperature controller purchased

from OMEGA instruments. The fittings used in the manifold are Swagelok ultratorr fittings. The

tubes are heated in a three-zone split furnace manufactured by Advanced Thermal Systems. It is

controlled by a Thermcraft three-zone temperature controller. The exhaust system consists of a

liquid nitrogen trap and a molecular sieve trap connected in series. The liquid nitrogen trap

condenses the unreacted precursor and prevents the vapors from getting into the pump. The

molecular sieve traps the organic products from the reactor. The pressure is measured by a

capacitance manometer and controlled by a butterfly throttle valve purchased from MKS

instruments. The exhaust gas is pumped out of the system with a rotary vane pump.

3.3 Experimental Procedure

The conditions used for the deposition of the coatings from the metalorganic precursors are

mentioned in Table 3.3. Alumina coating is deposited by MOCVD on tubes with diameters (Φ in

inches) 1/4 and 1/8 of an inch O.D. The tube to be coated is placed in a three zone split furnace

and heated to the reaction temperature TR. The tube is 30 cm long and the three-zone split

furnace helps to maintain the temperature constant throughout the length of the reactor. Alumina

has wide applications as membranes, insulating oxides and as good oxygen diffusion barrier [3]

.

MOCVD of alumina has been studied extensively from various precursors like trimethyl

aluminum [24,25]

, aluminum acetylacetonate [26-28]

, aluminum trisecondary butoxide [29-33]

and

aluminum isopropoxide [34,35]

. Two precursors aluminum acetylacetonate (Alacac) and aluminum

trisecondary butoxide (ATSB) were used for the deposition of alumina and the coatings obtained

from them will be designated as Alacac and ATSB respectively for convenience in the rest of the

77

thesis. The liquid precursor aluminum trisecondary butoxide is evaporated in the bubbler

whereas the solid precursor aluminum acetylacetonate is sublimed in the bubbler at the

temperature TB specified in Table 3.3. Platinum metal has been deposited on the surface of

alumina coating with the help of a solid precursor platinum acetylacetonate. The preheating lines

are heated to 190 °C to prevent the condensation of the precursor. The vapors are transported

through the preheating lines to the reactor with the help of carrier gas argon at a flow rate of Q

sccm and specified in Table 3.3. The reactor is maintained at a pressure of P Torrs and

mentioned in Table 3.3. Decomposition of the precursor on the surface of the tube results in the

formation of alumina. The unreacted precursor vapors leaving the reactor condense in the liquid

nitrogen trap and the product gases are pumped out of the reaction chamber. All the coating

experiments are performed for 3 hours unless specified.


3.4.1 Characterization of alumina coatings from aluminum trisecondary butoxide

Preliminary experiments were conducted to coat two fuel injectors for an aircraft engine.

Typical CVD process takes place at temperatures as high as 800 °C. To reduce the deposition

temperature to lower values, metalorganic precursors were preferred and their structures are

shown in Figure 3.3. Even though, the fluorinated metalorganic precursor aluminum

hexafluoroacetylacetonate shown in Figure 3.3a has a higher volatility, they are not preferred for

the deposition process, as the decomposition of the halogenated precursor results in the

formation of a corrosive hydrogen fluoride product. Therefore, a metalorganic precursor

aluminum isopropoxide, the structure of which is shown in Figure 3.3b, was used for the

deposition of alumina coating on the flow passages in the fuel injector. The injector to be coated

78

was placed in between two patterned aluminum blocks shown in Figure 3.4a. The modifications

made in the injector to facilitate the deposition of coatings by MOCVD are shown in Figure 3.4b

and Figure 3.4c. A check valve which is usually present in the fuel injector to ensure

unidirectional flow was removed to facilitate the coating process. The heating tape was wrapped

around the injector and heated to 550 °C. The carrier gas flow rate used in the experiment was

800 sccm. Ultra high purity oxygen at a flow rate of 100 sccm was used to ensure complete

decomposition of the metalorganic precursor. The system was maintained at a pressure of 20

Torr during the coating run. For the first experiment, the conduit shown in Figure 3.4b was used

as the inlet for CVD process. The tip of the injector shown in Figure 3.4c was used as an inlet for

the CVD process during the second experiment. An inspection of the fuel injector after the

deposition process provided guidance in choosing the parameters needed for subsequent tube

coating experiments. Aluminum isopropoxide was observed to be unstable during the deposition

process in the presence of oxygen. It is known to decompose into a less volatile species during

sublimation[35]

. The tip of the injector inspected after the deposition process (shown in Figure

3.4d) indicated the blockage of flow passages as small as 1/16‖ of an inch. Considering that, the

surface to volume ratio of a tube is inversely proportional to the diameter of the tube, the 1/16

inch tube has a higher surface-to-volume ratio compared to tubes of larger diameters present in

the flow passages of a fuel injector. Higher surface to volume ratio facilitates heterogeneous

reactions over homogeneous reactions[36]

. The higher rate of decomposition of the precursor on

the surface of 1/16‖ of an inch tube over a short duration of time may have blocked the flow

passages at the tip of the injector. It is also considered that higher carrier gas flow rates may have

carrier more precursor vapors. Therefore, the upper limit for the carrier gas flow rate was set at

400 sccm for subsequent experiments. The tip of the injector in Figure 3.4d showed the

79

formation of rust during the deposition process. Therefore, the upper limit for the deposition

temperature was set to 500 °C and oxygen was not used in the tube coating experiments.

The photographs of the uncoated and the coated tubes are shown in Figure 3.5a-c

respectively. The coating ATSB-A is deposited by the sublimation of aluminum trisecondary

butoxide at 160 °C and the subsequent pyrolytic decomposition of these vapors at 400 °C on a

¼‖ electrolposished stainless steel tube surface. The pressure inside the tube during the

deposition experiment is maintained at 10 Torrs. The conditions are chosen to avoid the

depletion of the precursor concentration along the length of the reactor in order to investigate the

possibility of coating a tube that is 30 cm long and evaluate the morphology of the coating. The

photograph of the coating ATSB-A is shown in Figure 3.5b. The morphology of the coating, the

EDX elemental mapping of aluminum, oxygen, carbon and the EDX spectrum are shown in

Figure 3.6.a-e respectively. The morphology of the coating in Figure 3.6a shows pores between

clusters. Within a cluster, there are faceted structures that come in contact with one another.

SEM measurements show that the average diameter of the cluster is 130 μm. The coating looks

brown in color which may be due to the presence of impurities like hydroxyl groups. The coating

was removed from the substrate and analyzed by X-ray diffractogram (XRD). Comparison of the

diffraction peaks shown in Figure 3.6f of the coating with the database in the International

Centre for Diffraction Data library (ICDD) shows the presence of boehmite. As the

decomposition temperature of the precursor is around 400 °C, dehydroxylation of the alumina

coating can be incomplete at these temperatures leading to the formation of boehmite that can be

brown in color. Investigation of alumina coating deposited at 400 °C by FTIR shows the

presence of boehmite [33]

. Cracks are present between the faceted structures. The EDX spectrum

shows the presence of aluminum, oxygen and almost no carbon. The higher intensity of

80

aluminum in the spectrum is an artifact introduced by gold, which shows a strong signal at the

same energy as aluminum, sputtered on the sample to avoid charging. The typical growth rate of

alumina from aluminum trisecondary butoxide varies between 0.35 μm/h and 1.4 μm/h [37]

. The

typical thickness of a TGO in a TBC varies between 25 – 125 μm [3]

. The cross-sectional SEM

micrograph in Figure 3.7a shows that the average thickness of the coating is 86 μm. Higher

bubbler temperature produces higher concentration of the precursor in the gas phase which in

turn increases the deposition rate and produces a thicker coating within 3 hours. The residual

stress of the coating increases with coating thickness [38]

. The photograph of this coating (Fig

3.5b) shows the delamination of the coating from the substrate surface as indicated by the red

circle. The delaminated region is examined by SEM. The morphology, elemental maps of

aluminum, oxygen and iron corresponding to the cross-section are shown in Figure 3.7a-d. The

blue rings drawn in the micrograph Figure 3.7a show the presence of pre-existing gaps that have

formed possibly during film growth between the coating and the substrate. The presence of these

interface defects is a prerequisite for spallation [38]

. Higher vapor pressure increases the

concentration of the precursor in the gas phase. This results in the formation of products by

homogeneous nucleation in the gas phase that might introduce these interfacial defects. The

spallation of oxide coatings is assumed to occur during cooling due to the thermomechanical

strain induced by the mismatch in the coefficient of thermal expansion between the substrate and

coating [39]

. As the coefficient of thermal expansion (CTE) of the substrate is greater than that of

the coating, the strain of the substrate exceeds the strain of the coating during the process of

cooling from the deposition temperature to room temperature. But the thermal stress induced due

to the mismatch in CTE is at less deposition temperatures as low as 400 °C compared to the

intrinsic stress. The tensile stress normal to the substrate develops during cooling and drives

81

spallation [40]

. The delamination of thick alumina coatings by interfacial cracking is observed [41]

.

It suggests the CVD process has produced a coating with a very high tensile residual stress. The

highest value of residual stress is induced at the interface as it is weaker than both the substrate

and the coating material [19]

. From elemental aluminum map, it appears that a thin alumina scale

may be adhering to the substrate after the spallation of the thick layer. Even though the coating

from aluminum trisecondary butoxide is thick and extends along the length of the tube, the

conditions chosen produced a porous morphology through which the fuel can access the

substrate surface to produce solid deposits and therefore make them unsuitable for the formation

of an environmental barrier coating along the entire length of the tube.

By decreasing the bubbler temperature, the concentration of the precursor in the gas phase

can be reduced to avoid homogeneous nucleation. This lowers the deposition rate. But,

eliminates cluster formation. The coating ATSB-B is deposited by the evaporation of the

precursor at 132 °C and its subsequent decomposition at 400 °C on the ¼’ electropolished

stainless steel tube surface. The photograph of the coating is shown in Figure 3.5c. The SEM

micrograph, the elemental map of aluminum, oxygen, carbon and the EDX spectrum of the

coating ATSB-B is shown in Figure 3.8a-e respectively. The micrograph of the coating in Figure

3.8a shows that the surface of the coating has a rough granular morphology. A closer look of the

coating at higher magnification shown in Figure 3.9a and Figure 3.9b reveals the presence of

cracks. The width of the crack in the coating ATSB-B is measured to be 330 nm as shown in

Figure 3.9b. At lower deposition temperature, the mobility of atoms on the surface of the

substrate is very low. This increases the intrinsic stress of the coating [32]

. The nature of intrinsic

stress in the CVD-grown alumina coating from the aluminum trisecondary butoxide was found to

be tensile at 400 °C and changes to compressive stress at 500 °C [28]

. At lower deposition

82

temperatures, intrinsic stress exceeds the thermal stress. Cracks on the surface are due to the

relaxation of tensile stress during the process of cooling the system from the deposition

temperature to the room temperature [42]

. The conditions mentioned for the formation of the

coating ATSB-B are therefore not suitable for the formation of an environmental barrier coating.

To reduce the intrinsic stress, improve the surface mobility of the atoms in the substrate,

adhesion and density of the coating, the deposition was carried out at a higher temperature. The

coating ATSB–C is deposited by evaporating aluminum trisecondary butoxide at 132 °C and the

subsequent decomposition of the precursor occurs at 450 °C on the ¼’ electropolished stainless

steel surface at a pressure of 10 Torr. A closer look at the microstructure of the coating in Figure

3.10a shows the presence of cauliflower morphology with pores. The cauliflower morphology

indicates that the deposition involves the formation of a solid by homogeneous gas phase

reaction [43,44]

. To reduce the gas phase reaction, increase the diffusion to the surface and thereby

improve adhesion of the coating, the pressure of the reactor is reduced and the temperature of the

substrate is increased in the subsequent experiment. By the evaporation of aluminum

trisecondary butoxide at 132 °C and subsequently decomposition of the vapors at 500 °C, the

coating ATSB-D is deposited at 500 °C on the inner surface of a ¼‖ electropolished stainless

steel tube at a reactor pressure of 2 Torr. The morphology of the coating ATSB-D is shown in

Figure 3.10b. The coating ATSB-D is also porous and has cauliflower morphology. The size of

the cluster is small and the void between cauliflowers is reduced in ATSB-C relative to that in

ATSB-D. Based on the EDX signal for aluminum and oxygen which is used to map the coating

along the length of the tube, at a deposition temperature of 500 °C, the deposition process began

at 8 cm from the inlet of the reactor and the length of the coated region was 7 cm. The hardness

of the coating with cauliflower morphology is lower than that with faceted morphology [44]

.

83

Aluminum in aluminum tri-secondary butoxide monomer has a coordination number of three.

The metal in the metal organic precursor always has a tendency to maximize its coordination

number by forming an alkoxide bridge in order to minimize its free energy [45]

. Decomposition of

the oligomer into the monomer increases the concentration of the reactant in the gas phase [46]

which would promote homogeneous reactions. Therefore, the film grown under this condition is

not suitable to be an environmental barrier to inhibit carbon deposition. The formation of pores

and the inability of the coating to protect the substrate against sulfidation under similar

conditions was reported in the literature [33]

. The conditions chosen for the deposition of ATSB-

A, ATSB-B, ATSB-C and ATSB-D from aluminum tri-secondary butoxide mentioned in Table

3.3 fall under mass-transport limited regime. It is reported that decomposition of alumina in the

mass-transport limited regime results in the formation of pores due to the rapid rate of

decomposition of aluminum tri-secondary butoxide which takes place by β-hydride elimination

reaction [46]

with an apparent activation energy value for the heterogeneous reaction reported to

be 65 kJ/mol [33]

. Further reduction in the reactor pressure did not improve the morphology of the

coating. The upper limit for the deposition temperature is set at 500 °C, because preliminary

coating experiments damaged the surface of a fuel injector at a deposition temperature of 550 °C.

The shelf life of the precursor appears to play a key role in determining the morphology of the

coatings in the tube obtained from aluminum trisecondary butoxide. Longer storage time in shelf

facilitates the oligomerization of the precursor. Subsequent rapid decomposition of the oligomer

at high temperature produces a porous coating. At lower deposition temperatures, the tubes can

be coated over the entire length. But the coating has cracks on the surface and suffers from

delamination. At higher deposition temperatures, the tubes can be coated to a length of 7 cm and

the coating is porous. Therefore, further improvements of the properties of the alumina coating

84

from aluminum trisecondary butoxide appears not to be possible due the chemical nature of the

precursor and the temperature limitation set by the process.

3.4.2 Characterization of alumina coatings from aluminum acetylacetonate

The vapor pressure of aluminum trisecondary butoxide at 134 °C is 0.8 Torr [31]

. The

vapor pressure of aluminum acetylacetonate at 138 °C is 0.23 Torr [47]

. Lower concentration of

the precursor in the gas phase reduces the homogeneous gas phase nucleation to a greater extent

that can avoid the deposition of clusters from the gas phase to the substrate surface and can

improve adhesion of the coating [38]

. Therefore, instead of aluminum trisecondary butoxide,

aluminum acetylacetonate is used as a precursor for the growth of alumina to improve the

morphology and adhesion of the environmental barrier coating. The coating AlacacA is

deposited by the sublimation of the aluminum acetylacetonate and the subsequent decomposition

of the precursor at 400 °C in a ¼‖ electropolished stainless steel tube surface at a reactor pressure

of 6 Torr.. The photograph of the coated tube is shown in Figure 3.11. The pink color of the

coating is mainly due to the interference of the light rays reflected between the film-air interface

and film-substrate interface. As the reactor is operated in the reaction rate controlled regime, the

thickness of the coating is uniform as seen from the uniformity of the pink color. The SEM

micrograph and EDX elemental map of aluminum, oxygen and carbon corresponding to AlacacA

are shown in Figure 3.13a-c respectively. The micrograph shows the presence of spherulitic

features which are the characteristic of the coating. Expansion of a gas containing a condensable

vapor like aluminum secondary butoxide through a subsonic nozzle in a high pressure system

was shown to produce condensate particles whose size distribution depends on the process

parameters like velocity of the fluid, reservoir temperature, pressure and saturation ratio [49]

. The

formation of spherulitic deposits has also been observed during the deposition of alumina from

85

aluminum acetylacetonate in a previous study and has been attributed to the condensation of

vapors into clusters during the expansion of a precursor-laden gas through a needle valve [50]

.

The system used in this study for the deposition of metal oxide films does not contain a needle

valve. Spherulitic features were not observed in the alumina coating from aluminum tri-

secondary butoxide discussed in the previous section. Therefore, the formation of spherulitic

features during the deposition of alumina from aluminum acetylacetonate appears to depend

upon the nature of ligand attached to the metal in the metal organic precursor. The EDX

elemental map in Figure 3.13d shows that carbon is homogeneously distributed throughout the

coating. The map also indicates that the spherulitic deposits contain both aluminum and oxygen.

TEM examination of the alumina coating deposited at 400 °C shows the presence of diffuse

diffraction rings indicating that the film deposited under these conditions is amorphous in nature.

Analysis on the coated tube surface by XPS showed the presence of aluminum, oxygen and

carbon. Usually, charging takes place when XPS is run on a non conducting ceramic substrate.

But in this sample, charging and a shift in the spectrum was not observed which is probably due

to the presence of a thin conductive coating. Depth profiling of the coatings suggested that the

coating thickness is around 30 nm.

In order to produce a thicker coating AlacacB, the deposition experiment was carried out

at a higher deposition temperature of 500 °C. The photographs of the uncoated and coated tubes

are shown in Figure 3.14 a-b respectively. The tube is coated between 10 cm and 15 cm from the

entrance of the reactor. At 500 °C, the reactor is operated in the mass transport limited regime.

Diffusion of the reactant across the boundary layer limits the rate of the process. The higher rate

of decomposition of the precursor at 500 °C causes a rapid depletion of the precursor

concentration in the gas phase. The SEM micrographs of the coatings and the respective

86

elemental maps of aluminum, oxygen and carbon are shown in Figure 3.15a-d. Comparing the

SEM micrograph of AlacacA in Figure 3.13a with that in Figure 3.15a, it can be observed that

the size of the spherulitic features has increased. The higher deposition temperature increases the

surface diffusion of all the spherulitic features that would agglomerate with one another and

increase in size. The thickness of the coating measured by depth profiling of XPS was found to

be around 500 nm. Heating the coated tube from 100 °C to 900 °C in the presence of ultra high

purity oxygen at a ramp rate of 30 °C/min and holding at 900 °C for 5 minutes did not result in

any spallation. Survey spectrum of XPS on the heated tube is shown in Figure 3.16a. The

spectrum shows the presence of O 1s, Al 2s and Al 2p photoelectron peaks. The SEM

micrograph in Figure 3.16b shows the presence of some pores in the coating. For a given amount

of residual stress, the interfacial stress is directly proportional to the thickness of the coating and

inversely proportional to the radius of curvature of the tube [19]

. As the thickness of the coating

AlacacB is small relative to the radius of curvature of the tube, the magnitude of interfacial stress

is much smaller than that needed to cause delamination of the coating. This shows a strong

adhesion of the coating to the substrate. The composition of AlacacB was obtained on four

different spots chosen randomly along the length of the tube. Comparing the photograph of

AlacacB in the Figure 3.14a with the composition values in Table 3.4, it can be seen that the

composition of the coating is nearly uniform along the length of the tube.

The surface to volume ratio of the tube increases with decreasing tube diameter. Based on

the coating experiments with the fuel injector, it is essential to evaluate the process to coat tubes

of small diameter. Therefore, a tube with an internal diameter one-eighth of an inch was used for

a coating experiment. The coating AlacacC is deposited by the sublimation of the precursor

aluminum acetylacetonate at 138 °C and its subsequent decomposition on the internal surface of

87

an one-eighth of an inch tube at 500 °C at a reactor pressure of 24 Torrs. The coating was

examined with XPS. The photograph, morphology and the EDX elemental map of the coating

AlacacC are shown in Figure 3.17a-e. The survey spectrum of the coating shown in Figure 3.17f

after a deposition time of three hours shows the presence of a background in the spectrum due to

the metal substrate suggesting the deposition of a thin coating. The morphology of the coating as

seen in the SEM micrograph Figure 3.17a is very smooth. In order to get a thick coating, the

experiment was carried out under the same conditions for 24 hours. The morphology of the

coating observed in Figure 3.18a shows the presence of spherulitic features. The formation of

spherulitic features on the alumina coating from aluminum acetylacetonate is discussed in the

beginning of this section. In the system under study, the gas flows from a preheating line of a ¼‖

diameter to the 1/8‖ tube to be coated that are connected together with a Swagelok adapter. In

this situation the gas can expand only due to the pressure gradient caused by the pump in the

system. This contrasts with the observation of spherical features during the expansion of a gas

containing a condensable vapor like aluminum trisecondary butoxide from a high pressure region

to the low pressure region through a nozzle due to condensation. Therefore, the formation of

spherulitic features appears to be associated with the decomposition of the precursor containing

the acetylacetonate ligand. The EDX spectrum of the coating in Figure 3.18b shows a strong

signal for aluminum and oxygen. The signal from the substrate is not observed in the

background. This suggests that the deposition experiment conducted for a longer duration of time

produces a thicker coating on the shorter tube. The conductance of a tube is defined as the

amount of gas that can be pumped out of the conduit per unit time and is directly proportional to

the third power of the diameter and inversely proportional to the length of the tube. The diameter

and hence the conductance of the tube to be coated is lower than that of the preheating line and

88

the bubbler. The overall conductance of the system decreases when a tube with lower

conductance is connected in series with a preheating and a bubbler with higher conductance.

Therefore, due to the impedence of the small diameter tube with lower conductance, longer

duration of time is needed to reduce the pressure in the preheating line and the bubbler to lower

values and enhances the volatility of the precursor. The higher volatility increases the

concentration of the precursor in the gas phase which would increase the coating thickness. The

coatings from aluminum acetylacetonate are smooth and non-porous. Cracking and delamination

were not observed in the alumina coating obtained from aluminum acetylacetonate. Therefore

aluminum acetylacetonate appears to be a good candidate for the deposition of an environmental

barrier coating on complex geometries.

3.4.3 Temperature-Programmed Oxidation of Residual Carbon in Alumina Coatings

XPS results and EDX analysis suggest that the coating obtained from aluminum

acetylacetonate has a high concentration of residual carbon due to the incomplete decomposition

of the precursor. The formation of an intermediate compound aluminum hydroxyl

acetylacetonate in the gas phase has been postulated which subsequently decomposes on the

surface to produce alumina [51]

. The apparent activation energy for the pyrolysis of the precursor

is reported to be 108 kJ / mol. The amount of residual carbon in the alumina film obtained from

the pyrolysis of the precursor aluminum acetylacetonate has been measured by temperature

programmed oxidation. The alumina coating was deposited at the substrate temperatures of 350 °

C, 375 ° C, 400 ° C and 450 ° C on AISI 304 for a period of 24 hours. The deposition time was

increased significantly in order to get a good signal to noise ratio for carbon dioxide in the

RC412 multiphase carbon analyzer. The length of each sample is 5 cm. Each sample is loaded in

the quartz boat and heated from 100 °C to 900 °C at a ramp rate of 30 °C/min in the presence of

89

UHP oxygen and held at 900 °C for 5 minutes. Figure 3.19a shows the TPO profiles of the

residual carbon for the respective samples. The amount of residual carbon decreases with the

increasing deposition temperature due to the more complete decomposition of the precursor. The

low temperature peak in all the samples is due to the presence of structurally disordered reactive

carbon that oxidizes at relatively lower temperatures between 260 °C and 400 °C. The peak

between 400 °C and 600 °C in all the samples suggest the oxidation of relatively less reactive

hydrogen lean structurally more ordered carbon. The high resolution scan for the carbon 1s

shown in the Figure 3.19b present in the alumina coating deposited at 450 °C has been

deconvoluted by the CASAXPS software. Deconvolution of the peak shows three components.

By comparison of these peaks to the standard values for the polymer poly(acetylacetoxyethyl

methacrylate) in the XPS reference handbook, the first peak with a binding energy maximum at

284.6 eV is assigned to aliphatic carbon (C-C, C-H), the second peak with a binding energy

maximum at 286 eV is assigned to the presence of –CH2 group present between carboxyl groups

and the third peak with a binding energy maximum at 288 eV is assigned to the presence of

carboxyl groups.

3.4.4 Characterization of Platinum Deposited on Alumina Coatings

Platinum metal was deposited on the surface of alumina coating obtained from aluminum

acetylacetonate AlacacB. The metal was deposited by the sublimation of platinum

acetylacetonate at 138 °C and the subsequent decomposition of the precursor at 500 °C. The

reactor is maintained at a pressure of 6 Torr. The deposition process was conducted subsequent

to alumina deposition from aluminum acetylacetonate and the reactor was not cooled down

during the process of switching the precursor in the bubbler. The photograph of the coating is

shown in Figure 3.20. The morphology of the coating shown in Figure 3.21a looks very much

90

like that of AlacacC shown in Figure 3.18a. The surface composition of the coatings AlacacB

and Pt-AlacacB was found by acquiring the high resolution scan for the elements Al 2p, O 1s, C

1s and Pt 4f in X-ray photoelectron spectroscopy. The results are tabulated in Table 3.4.

Comparison of the surface composition of the elements in these coatings shows that the

concentration of platinum in atomic percentage deposited on the alumina coating from aluminum

acetylacetonate denoted as Pt-AlacacB-1hr increases from 0.6% for a coating duration of 1 hour

to 2.3% for the coating Pt-AlacacB-3hr that is deposited for 3 hours. The high resolution scan for

Pt 4f is shown in Figure 3.21b. The peak height on the right side corresponding to 4f7/2 is shorter

than that corresponding to 4f5/2 for the sample whereas the peak height for 4f7/2 should be greater

than that for 4f5/2 for metallic platinum. Deconvolution of the high resolution scan for platinum

4f shows two components. The position of the first component with a binding energy

corresponding to platinum 4f7/2 at 71.3 eV suggests that platinum is present in the form of a

metal. This is represented by the blue line in Figure 3.21b. The position of the second component

represented by the pink line with a binding energy corresponding to 74.2 eV suggests that

platinum is also present in the form of metal oxide with an oxidation state of +2.

The higher content of carbon in the coating Pt-AlacacB may be due to the incomplete

decomposition of the platinum precursor. The process of sectioning the tubes after coating using

a grinding wheel may also have contributed significantly to the presence of adventitious carbon

even though the tubes were rinsed with acetone and ethanol after the sectioning of the tube. The

deconvolution of the high resolution scan of C 1s into its individual components shows the

presence of C-C (285 eV), C-O (287 eV) and COO (289 eV) and gives their relative

concentration on the coating. These values are tabulated in Table 3.5. Using the concentration of

hydrocarbon species present, the thickness of the organic overlayer of these species has been

91

predicted using modified form of Smith’s correction formula [52]

. The concentration of the

functional groups in the organic overlayer is less for alumina coating before platinum deposition

and increases with increase in the deposition time for platinum coating suggesting the

incorporation of residual carbon impurities during the deposition of platinum at 450 °C in the

absence of oxygen.

3.5 Conclusions

Thick alumina coatings deposited from aluminum trisecondary butoxide at 400 °C suffer

from cracks and pre-existing separation from the substrate surface due to the intrinsic stress in

the coating. The pre-existing separation between the coating and the substrate suggests the

presence of normal stress component perpendicular to the film-substrate interface that reduces

the resistance of the coating to spallation. The formation of surface cracks on the coating

suggests the relaxation of the intrinsic tensile stress in the coating due to the lower deposition

temperature. Oligomerization of aluminum trisecondary butoxide during storage and its

subsequent decomposition during the deposition process increases the concentration of the

reactant in the gas phase and produces a porous coating at 500 °C. The shelf life of the precursor

facilitates oligomerization and affects the reliability of the precursor. Therefore, the coatings

produced from aluminum trisecondary butoxide are not good environmental barriers.

Aluminum acetylacetonate has a lower vapor pressure compared to aluminum

acetylacetonate and the alumina coating, deposited at 400 °C is thin, free from cracks and

uniform in thickness as a function of reactor length. The formation of the spherulitic features in

the coating appears to be associated with the decomposition of the acetylacetonate ligand. A

higher deposition temperature of 500 °C increases the surface diffusion and leads to the

92

agglomeration of these spherulitic features on the substrate surface. The morphology shows that

the alumina coating from aluminum acetylacetonate is smoother than that from aluminum

trisecondary butoxide. At higher decomposition temperature, the concentration of the precursor

decreases rapidly in the gas phase resulting in the non-uniformity in the thickness of the coating.

But the composition of the coating appears to remain uniform over the coated length. The

coating contains residual carbon left due to the incomplete decomposition of the precursor.

Deconvolution of high resolution scan for C 1s in the coating shows the presence of carboxyl

groups in the coating. Heating the 500 nm thick coating from 100 °C to 900 °C at a ramp rate of

30 °C/min in the presence of UHP oxygen does not cause delamination. The thinner coating in a

tube of relatively large radius of curvature produces less normal stress that may not cause

delamination of the coating. The amount of residual carbon incorporated in the coating decreases

with increasing deposition temperature. TPO indicates the presence of structurally disordered

hydrogen-rich carbon oxidizing at lower temperatures between 260 °C and 400 °C and relatively

less reactive structurally more ordered carbonaceous species oxidizing at temperatures between

400 °C and 600 °C. Platinum deposited from platinum acetylacetonate on the alumina coating

Alacac has two oxidation states one corresponding to the metal and the other corresponding to

the metal oxide with an oxidation state of +2. The amount of platinum deposited on the alumina

coating increase from 0.6 atomic percentage to 2.3 atomic percentage when the deposition time

increases from 1 to 3 hours.

93

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31. Haanappel, V. A. C.; Van Corbach, H. D.; Fransen, T.; Gellings, P. J., The effect of thermal

annealing on the properties of alumina films prepared by metal organic chemical vapour

deposition at atmospheric pressure. Surface and Coatings Technology 1994, 64 (3), 183-193.


films prepared by low-pressure metal-organic chemical vapour deposition. Surface and Coatings

Technology 1995, 72 (1-2), 13-22.

33. Haanappel, V. A. C.; Vendel, D. V. D.; Metselaar, H. S. C.; Van Corbach, H. D.; Fransen, T.;

Gellings, P. J., The mechanical properties of thin alumina films deposited by metal-organic

chemical vapour deposition. Thin Solid Films 1995, 254 (1-2), 153-163.

34. Aboaf, J. A., Deposition and properties of aluminum oxide obtained by the pyrolytic

decomposition of aluminum alkoxide. Journal of the electrochemical society 1967, 114 (9), 948-

952.

35. Gleizes, A. N.; Vahlas, C.; Sovar, MM.; Samélor, D.; Lafont, MC., CVD-Fabricated

aluminum oxide coatings from aluminum tri-iso-propoxide: Correlation between processing

conditions and composition. Chemical Vapor Deposition 2007, 13, 23-29.

36. Huttinger, K. J.; CVD in Hot Wall Reactors—The Interaction Between Homogeneous Gas-

Phase and Heterogeneous Surface Reactions. Chemical Vapor Deposition 1998, 4(4), 151-158.

37. Kuo, D. H.; Chuang, P. Y., Aluminum Silicate Films Obtained by Low-Pressure Metal-

Organic Chemical Vapor Deposition. J. Am. Ceram. Soc. 2003, 86 (6), 969-974.

38. Evans, A. G.; Crumley, B. C.; Demaray, R. E., On the mechanical behavior of brittle

coatings and layers. 1983, 20 (5-6), 193-216.

39. Harvey, M.; Courcier, C.; Maurel, V.; Rémy, L., Oxide and TBC spallation in [beta]-NiAl

coated systems under mechanical loading. Surface and Coatings Technology 2008, 203 (5-7),

432-436.

40. Hou, P. Y.; Tolpygo, V. K., Examination of the platinum effect on the oxidation behavior of

nickel-aluminide coatings. Surface and Coatings Technology 2007, 202 (4-7), 623-627.

41. Haanappel, V. A. C.; Van Corbach, H. D.; Fransen, T.; Gellings, P. J., Cracking and

delamination of metal organic vapour deposited alumina and silica films. Materials Science and

Engineering A 1993, 167, 179-185.

96

42. Ashwin R. Shah; David N. Brewer; Pappu L. N. Murthy, Life Prediction Issues in

Thermal/Environmental Barrier Coatings in Ceramic Matrix Composites. 2001.

43. Zayan, M. H., Model for Nonprotective Oxidation of Al-Mg Alloys. Oxidation of Metals

1990, 34 (5-6), 465-472.

44. Konoplyuk, S.; Abe, T.; Takagi, T.; Uchimoto, T., Hot filament CVD diamond coating of

TiC sliders. Diamond and Related Materials 2007, 16 (3), 609-615.

45. Bradley, D. C., Metal Alkoxides as Precursors for Electronic and Ceramic Materials.

Chemical Reviews 1989, 89, 1317-1322.

46. Haanappel, V. A. C.; Van Corbach, H. D.; Fransen, T.; Gellings, P. J., The pyrolytic

decomposition of aluminium-tri-sec-butoxide during chemical vapour deposition of thin alumina

films. Thermochimica Acta 1994, 240, 67-77.

47. Pflitsch, C.; Muhsin, A.; Bergmann, U.; Atakan, B., Growth of thin aluminium oxide films

on stainless steel by MOCVD at ambient pressure and by using a hot-wall CVD-setup. Surface

and Coatings Technology 2006, 201 (1-2), 73-81.

48. Choy, K. L., Chemical vapor deposition of coatings. Progress in Materials Science 2003, 48,

57-170.

49. Turner, J. R.; Kodas, T. T.; Friedlander, S. K., Monodisperse particle-production by vapor

condensation in nozzles. J. Chem. Phys. 1988, 88 (1), 457-465.

50. Singh, M. P.; Shivashankar, S. A., Low-pressure MOCVD of Al2O3 films using aluminium

acetylacetonate as precursor: nucleation and growth. Surface and Coatings Technology 2002,

161 (2-3), 135-143.

51. Rhoten, M. C.; DeVore, T. C., Evolved Gas Analysis Investigation of the Reaction between

Tris(2,4-pentanedionato)aluminum and Water Vapor in Chemical Vapor Deposition Processes

To Produce Alumina. Chemistry of Materials 1997, 9 (8), 1757-1764.

52. Smith, G. C., Evaluation of a simple correction for the hydrocarbon contamination layer in

quantitative surface analysis by XPS. Journal of Electron Spectroscopy and Related Phenomena

2005, 148 (1), 21-28.

97

Table 3.1. Benefits and limitations of various deposition methods [1]

Features

Process

Evaporation Sputtering CVD Electro-deposition Thermal Spray

Mechanism

to produce

species

Heat Momentum Chemical

Reaction

Solution Plasma or

flames

Deposition

rate

Moderate Low Moderate Low to high Very high

Deposition

species

Atoms Atoms/ions Atoms/ions Ions Droplets

Complex

shapes

Poor line of

sight

Good but

not uniform

Good Good Poor resolution

Deposits in

small blind

holes

Poor Poor Limited Limited Very limited

Metal/alloy

deposition

Yes Yes Yes Yes Yes

Ceramic

deposition

Yes Yes Yes limited Yes

Energy of

deposit

species

low can be high can be high can be high can be high

Growth

interface

perturbation

Not

normally

Yes Yes No No

Substrate

heating

Yes Not

normally

Yes No Not normally

98

Table 3.2. Description of the components used in the MOCVD system

Table 3.3. Growth conditions used for alumina coating from aluminum trisecondary

butoxide (ATSB) and aluminum acetylacetonate (Alacac)

Serial no Components Model No

1

2

3

4

5

6

7

8

9

10

11

12

13

Mass flow controller for carrier gas argon

Mass flow controller for purge gas argon

Mass flow controller for UHP oxygen

Bellow Sealed valve for open and close applications

Bubbler for precursor evaporation and sublimation

Bellow sealed valve to exclude bubbler for leak testing

Three-zone temperature controller for the reactor

Three-zone split furnace

Capacitance manometer

Liquid nitrogen trap

Molecular sieve trap

Butterfly control valve

Rotary vane pump

FMA-765A-V02.5/0 psi



SS-4BG-V51

Custom-built

SS-4BG-V51

3D1-36-208

115 V 4.5 A/9 A

627B13TBC1B-MKS

TNR4XR075QF-BOC

TSR4MS075QF-BOC

253B-11020-MKS

RV8-BOC Edwards

Sample Φ TB° C Q (sccm) TR ° C P (Torr)

ATSB-A

ATSB-B

ATSB-C

ATSB-D

¼

¼

¼

¼

160

132

132

132

400

400

400

400

400

400

450

500

10.0

10.0

10.0

2.0

AlacacA

AlacacB

AlacacC

Pt-AlacacB

¼

¼

1/8

1/4

138

138

138

138

100

100

100

100

400

500

500

400

6.0

6.0

24.0

6.0

99

Table 3.4. Composition of the coating by X-ray photoelectron spectroscopy

Relative concentration of elements in atomic percentage of the coating from high resolution scan

by XPS

Sample Al Pt O C

AlacacB-site 1 27.0 0.0 45.2 27.8

AlacacB-site 2 28.2 0.0 45.9 25.9

AlacacB-site 3 28.9 0.0 48.4 22.7

AlacacB-site 4 28.1 0.0 44.9 24.4

Pt-on AlacacB-1hr 20.6 0.6 44.0 34.8

Pt-on AlacacB-3hr 11.4 2.3 24.9 61.4

Table 3.5. Concentration of various functional groups under C 1s peak in atomic percentage

Sample C-C C-O COO Overlayer thickness Ǻ

AlacacB-site 1 20.8 3.3 3.7 15.2

AlacacB-site 2 18.6 3.9 3.3 14.3

AlacacB-site 3 16.5 3.0 3.2 12.1

AlacacB-site 4 16.7 5.1 2.5 13.7

Pt-on AlacacB-1hr 26.8 3.2 4.8 19.9

Pt-on AlacacB-3hr 49..8 4.6 3.2 41.7

100

Figure 3.2. Schematic representation of the MOCVD system for coating tubes.

Figure 3.3. Structure of the metalorganic precursors.

101

Figure 3.4. (a) Photograph of the aluminum block used for heating the fuel injector (b)

Photographs of a fuel injector nozzle (c) Photographs of a fuel injector before and after the

attachment of the fitting at the tip (d) Photograph showing the injector tip before (right) and after

deposition (left).

102

Figure 3.5. (a) Photograph of an uncoated ¼‖electropolished stainless steel tube heated to 500

°C in the presence of argon (b) Photograph of the coating ATSB-A at 400 °C sectioned into two

halves (c) Photograph of the coating ATSB-B deposited at 450 °C.

Delamination of the coating from the substrate

Alumina coating

Electro polished ¼‖ stainless steel tube substrate

a

b

C

103

Figure 3.6. (a) SEM micrograph of alumina coating ATSB-A deposited from aluminum

trisecondary butoxide at TS = 400 °C and TB = 160 °C. EDX elemental map of (b) aluminum and

(c) oxygen (d) carbon (e) EDX signal corresponding to an elemental composition of 39% of

aluminum and 61% of oxygen (f) X-ray diffractogram of the coating ATSB-A scraped from the

substrate.

a c b

a

d e

Al

O Au

Al O

C

f

104

Figure 3.7. (a) Cross-sectional SEM micrograph of alumina coating ATSB-A deposited from

aluminum trisecondary butoxide at TS = 400 °C and TB = 160 °C. EDX elemental map of (b)

aluminum (c) oxygen (d) Iron.

a b c

d

Al O

Fe

105

Figure 3.8. (a) SEM micrograph of alumina coating ATSB-B deposited from aluminum

trisecondary butoxide at TR = 400 °C and TB = 132 °C. Elemental map of (b) Aluminum (c)

oxygen (d) carbon on the coating (e) EDX spectrum of the coating.

a b

c d

e

Al

O C

106

Figure 3.9. SEM micrographs of alumina coating ATSB-B deposited from aluminum

trisecondary butoxide at TS = 400 °C and TB = 132 °C.

Figure 3.10. SEM micrograph of alumina coating from aluminum trisecondary butoxide

evaporated at 132 °C and deposited at (a) 450 °C for ATSB-C (b) 500 °C for ATSB-D.

Figure 3.11. (a) Photograph of the coated tube AlacacA

a b

a

a b

107

Figure 3.12. (a) SEM micrograph of alumina coating AlacacA from aluminum acetylacetonate

sublimed at 138 °C and deposited at 400 °C (b) Elemental map of Aluminum (c) oxygen (d)

carbon (e) EDX spectrum of the coating.

a b

c d

Al

O C

108

Figure 3.13. (a) Photograph of an uncoated tube. (b) Photograph of the coated tube AlacacB

Figure 3.14. (a) SEM micrograph of alumina coating AlacacB from aluminum acetylacetonate

sublimed at 138 °C and deposited at 500 °C. (b) Elemental map of Aluminum (c) oxygen (d)

carbon.

a b

Al

O C

a b

d c d

109

Figure 3.15. (a) XPS survey scan of thermally stressed alumina film AlacacB on AISI 304 after

temperature programmed oxidation from 100 °C – 900 °C and holding the coating at 900 °C for

5 minutes in UHP oxygen.(b) SEM micrograph of the coating at 30μm

a b

b

e d

c

Al

O C

a

110

Figure 3.16. (a) Photograph of coated 1/8‖ tube (b) SEM micrograph of alumina coating

AlacacC from aluminum acetylacetonate sublimed at 138 °C and deposited at 500 °C after 4

hours at 30x. (c) Elemental map of Aluminum (d) oxygen and (f) carbon and XPS survey scan of

Alumina coating on the one eighth of an inch tube.

Figure 3.17. (a) SEM micrograph of alumina coating AlacacC from aluminum acetylacetonate

sublimed at 138 °C and deposited at 500 °C after 24 hours. (b) EDX spectrum of alumina coating

on the one eighth of an inch tube showing the presence of aluminum and oxygen.

f

b a

O

Al

111

Figure 3.18. (a) TPO of residual carbon in the alumina coating formed by the decomposition of

aluminum acetylacetonate. (b) High resolution scan of C 1s by XPS on the alumina coating

deposited at 450 °C.

Figure 3.19. Photograph of coated tube Pt-AlacacB

Figure 3.20. (a) SEM micrograph of the coating Pt-AlacacB (b) High resolution scan for

platinum 4f from X-ray photoelectron spectroscopy.

a

a b

Binding energy between

288-289 eV indicates the

presence of carboxyl groups

112

Chapter 4

Effectiveness of Low-Pressure MOCVD Coatings on Metal Surfaces for the Mitigation of

Fouling from Heated Jet Fuel

4.1 Abstract

Thin films of alumina, zirconia, tantalum oxide and platinum were deposited on AISI304

by metal organic chemical vapor deposition (MOCVD) to investigate the effectiveness of these

coatings in inhibiting carbon deposition and sulfide formation from thermal oxidative

degradation of jet fuel. Coated AISI304 foils were heated in a laboratory scale flow reactor with

a commercial jet fuel (Jet-A) flowing at 1 mL/min at a wall temperature of 350 °C and reactor

pressure of 500 psig (3.4 MPa) for 5 hours. Under these conditions, both liquid phase

autoxidation and thermal decomposition of jet fuel contribute to carbon deposition. The surface

composition of the metal oxide coatings was found by X-ray photoelectron spectroscopy (XPS).

The morphology of the coating and the carbonaceous deposits formed during thermal stressing

were examined by field emission scanning electron microscopy (FESEM). The amount of solid

carbonaceous deposits on the coated and uncoated surfaces was measured by temperature-

programmed oxidation (TPO). The effectiveness of the coatings in mitigating carbon deposition

was found to decrease in the following order Platinum > Ta2O5 > alumina from acetyl acetonate

> ZrO2 > alumina from aluminum trisecondary butoxide > AISI304. The coatings cover the

metal surface by forming a protective layer that inhibits the formation of metal sulfides from the

reaction of sulfur compounds in jet fuel with iron and nickel on stainless steel surfaces. The

variation in the activity of the coatings can be attributed to the interaction of oxygenated

intermediates formed by autoxidation during thermal stressing with coating surfaces having

different degrees of acidity.

113

4.2 Introduction

Carbonaceous solid deposits form on various metal surfaces in the fuel handling system of

aircraft gas turbine engines when hydrocarbons flow through them at high temperatures and high

pressures [1]

. The formation of solid deposits is of major concern in aircraft engines as it can plug

the filters and accumulate on valves, flow lines and fuel injector nozzles affecting the fuel system

operation and cause severe degradation of metal surfaces [2]

. The varying activity of metals

towards solid deposit formation from stressing of Jet-A at 350 °C and 500 psig is attributed to

their reaction with reactive sulfur compounds in the fuel and their interaction with oxygenated

intermediates formed by autoxidation under thermal oxidative conditions [3]

. Passivation of metal

surfaces with an inert metal oxide coating appears to delay the amount of deposits formed over a

period of time under thermal oxidative [4]

and pyrolytic conditions [5,6]

. Thermal stressing

experiments with JP-8 under pyrolytic conditions on inert surfaces suggest that the glass-lined

stainless steel surface is least active towards carbon deposition, even lower than that observed on

vapor-deposited silicosteel surfaces [6]

. Alumina protects metal surfaces against corrosion and

oxidation at high temperatures [7]

. Its high chemical stability, dielectric property, radiation

resistance and low permeability to alkali ions has led to an extensive literature about the film

growth of alumina from aluminum trisecondary butoxide (ATSB) [8]

, aluminum acetylacetonate

(Alacac) [9]

, aluminum chloride [10]

, and trimethyl aluminum [11]

by metal-organic chemical

vapor deposition (MOCVD). Zirconia films deposited by MOCVD have been examined as a

protective coating to prevent the oxidation of graphite [12]

and as an ionic conductor for solid

oxide electrolytes [13]

. The dielectric property of zirconia and tantalum oxide has led to the

investigation of chemical vapor deposition of Zirconia from zirconium tertiary butoxide along

with other oxides [14]

and tantalum oxide from tantalum pentaethoxide [15]

. The objective of this

114

paper is to investigate the effectiveness of platinum and four metal oxide coatings that include,

alumina deposited from the metal organic precursor aluminum trisecondary butoxide referred as

ATSB, alumina from aluminum acetylacetonate referred as Alacac, zirconia deposited from

zirconium acetylacetonate, and tantalum oxide deposited from tantalum pentaethoxide for

inhibiting the carbon deposition and metal sulfide formation during the thermal stressing of Jet-A

at 350 °C and 500 psig. The metallic platinum coating was deposited from platinum

acetylacetonate. It should be noted that this study has focused on comparing the performance of

different MOCVD-based metal oxide and metal coatings produced under selected conditions

without an effort for optimizing the deposition conditions for each coating.


4.3.1 MOCVD Experimental Set-up for Foil Coatings

Figure 4.1 shows the schematic diagram of the experimental setup used for depositing the

coatings. AISI 304 sheet purchased from Goodfellow Ltd is cut into foils that are 80 mm long,

3mm wide and 0.9 mm thick, washed in soap solution, rinsed with acetone and ethanol and dried

in vacuum for an hour before the MOCVD experiment. The foil substrate is inserted inside a

stainless tube reactor that is 6.35 mm (1/4 inch) in diameter and 300 mm long. The edges of the

foil are fitted tightly between the walls of the reactor so that heat is conducted from the reactor

wall to maintain a constant temperature on foil surfaces. The metalorganic precursors used for

the coating are placed in bubblers that are heated to the respective sublimation temperatures. The

conditions used for the deposition of four coatings by MOCVD are summarized in Table 4.1.

Before each deposition, the system is evacuated to a pressure of less than 0.1 Torr and the reactor

tube is then heated to the reaction temperature (500 °C) in a three-zone split furnace in the

115

presence of UHP argon. The coatings deposited at temperatures less than 400 °C had significant

amounts of carbon. Reheating these coatings during the thermal stressing test discussed below

makes the coating porous and deteriorates the adhesion of the coating on the substrate.

Therefore, the substrate temperature was fixed at 500 °C so that the coating adheres to the

substrate well and does not deteriorate upon heating to 350 °C for thermal stressing experiments.

The sublimation temperatures of liquid precursor aluminum trisecondary butoxide and tantalum

pentaethoxide used in the deposition of alumina ATSB and tantalum oxide are 132 °C [8]

and 110

°C [16]

, respectively. The solid precursor aluminum acetylacetonate, zirconium acetylacetonate,

and platinum acetylacetonate used in the deposition of alumina Alacac, zirconia and platinum are

sublimed at 138 °C [17]

, 160 °C [18]

, and 145 °C [19]

, respectively. The gas flow rate, pressure and

the vaporization or sublimation temperatures in the bubbler are chosen such that a smooth and

uniform film is deposited over a 5 cm length of the foils. The precursor vapors are transported

into the reactor by ultra high purity argon through the preheating lines which are maintained at a

temperature 10 °C higher than the respective sublimation temperatures. Ultra high purity oxygen

is used as an oxidant for the decomposition of aluminum acetylacetonate and zirconium

acetylacetonate, as pyrolysis of the precursors leave much carbon on the coating that affects the

adhesion of the coating during the subsequent thermal stressing experiment. Once the reactor

reaches the reaction temperature, the carrier gas is allowed to flow through the bubbler instead of

the purge gas. The experiment is conducted for a period of four hours. At the end of the

experiment, the carrier and the oxidant gases are shut off and the purge argon is flown to cool the

reactor. Once the reactor reaches the ambient temperature, the system is brought back to

atmospheric pressure.

116


Analysis of Fuel: A commercial Jet-A fuel sample used in the thermal stressing was

analyzed by Shimadzu GC-17A gas chromatograph (GC) equipped with a QP-5000 Shimadzu

mass spectrometer (MS) for hydrocarbons. The column is 30 m long and 0.25 mm diameter. The

stationary phase used in the column is 5% phenyl-95% methyl polysiloxane. Helium is used as a

carrier gas. A 1 μl of sample is injected into the injector at 290 °C. The column is maintained at

40 °C for 4 minutes, and then heated to 220 °C at a ramp rate of 4 °C per minute and held at that

temperature for 10 minutes. To analyze the sulfur compounds in Jet-A, GC-PFPD (pulsed flame

photometric detector) HP5890 Series II was used. A 0.1 μl of the sample is injected into the

injector at 290 °C. The column was maintained at 120 °C for 2 minutes, ramped to 170 °C at the

rate of 6 °C per minute and subsequently to 290 °C at the rate of 20 °C per minute and held at

290 °C for 5 minutes. The total run time for the sample is 21.33 minutes.

Stressing Experiments: The uncoated and the coated substrate samples are washed in hexane

and dried in argon for an hour before each experiment. The experimental setup for thermal

stressing of Jet-A is shown in Figure 4.2 [3]

. The details of the thermal stressing reactor used for

isothermal experiments are described elsewhere [2]

. Stressing experiments are conducted in a

6.35 mm diameter (1/4-in o.d.), 20 cm long glass-coated stainless steel reactor. The 10-cm long

substrates are inserted to stay at the bottom section of the reactor. The reactor with the substrate

is heated in the presence of argon to 350 °C at a reactor pressure of 500 psig (3.5 MPa) with the

help of a block heater to maintain isothermal conditions along the length of the reactor and

maintained at that temperature for 4 hours to reach thermal equilibrium. Ultra zero air is bubbled

into the Jet-A reservoir so that it is saturated with dissolved oxygen during the course of the

experiment. The fuel is pumped into the system at 500 psig. It enters the preheating line of 3.175

117

mm diameter (1/8-in o.d.) and 2 m in length. The residence time of the fuel in the preheating line

is 6.3 minutes. It is preheated to 260 °C before entering the reactor. The fuel flow rate, reactor

wall temperature and the pressure are maintained at 1 mL/min, 350 °C, and 500 psig for 5 h. The

residence time of the fuel in the reactor is 1.4 minutes. At the end of the experiment, the residual

fuel in the reactor is removed by purging it with argon.

Preliminary experiments were conducted with AISI 316 and AISI 304 foils to study the

variation in deposit formation as a function of reactor length. The length of the foils used in this

experiment was 10 cm. The foils are located between 7.5 cm and 17.5 cm from the top of the

reactor [2]

. Temperature-Programmed Oxidation (TPO described below) conducted on the two

sections of the same substrate material, each 5 cm long, has shown the same amount of deposits.

The nature of the TPO curves corresponding to each substrate is discussed in the next section.

The TPO curves from each of the two sections of the same substrate were similar. Therefore, for

the TPO of other substrates, the portion of the foil located between 10 and 15 cm from the top of

the reactor was chosen for analysis. The total amount of solid carbonaceous deposits on the blank

substrate and the coatings is reproducible to within 5% of the deposit mass.

4.3.3 Characterization of Coatings and Carbon Deposits

FESEM and TEM Microscopy: The morphologies of the coatings after MOCVD, the

carbonaceous deposits, and coatings after thermal stressing are characterized by field emission

scanning electron microscopy JEOL 6700F (FESEM). The cross section of the coated specimen

alumina ATSB, alumina Alacac and zirconia were prepared by FEI quanta focused ion beam

(FIB) for TEM characterization. For this analysis, one stub with the coated substrate and the

other with the TEM copper grid are loaded into the chamber adjacent to one another. The

118

instrument is equipped with a tungsten omniprobe manipulator which can be moved back and

forth between the stubs with the XYZ controls of the stage. Depositing platinum on the coated

substrate as a protective layer to prevent ion beam induced damage to the coating is the first step.

The micrograph in Figure 4.3a shows the presence of platinum strips marked as Pt deposited

initially by electrons and subsequently by ions. Milling the surface with gallium ions to cut a

wedge shaped cross-sectional specimen marked as W in Figure 4.3a from the coated sample is

the second step. The sample is milled such that one of the corners of the specimen was attached

to the coated substrate. The other edges and corners are detached from the substrate. This is

shown in Figure 4.3b. In the third step, the tungsten omniprobe manipulator in the instrument is

navigated in the chamber over the sample stub such that the tip of the probe is in physical contact

with the corner of the specimen that is separated from the sample during milling, marked as 1 in

Figure 4.3b. Fastening the omniprobe tip to the specimen by depositing platinum in and around

the region of contact is the fourth step. Finally milling the corner of the specimen connected to

the substrate, marked as 2 in Figure 4.3b, with gallium ions detached the specimen from the

substrate so that the omniprobe fastened with the specimen could move freely towards the TEM

grid and put the other end of the specimen in contact with the TEM copper grid. Platinum

deposition in and around the region of contact fastens the specimen to the grid. The probe is

disconnected from the other end of the specimen by milling. The specimen bonded to the TEM

grid is shown in Figure 4.3c. The cross-section obtained is 10 μm long, 3 μm deep and 500 nm

wide. Milling the cross-section with electrons reduces the width of the sample and makes it

electron transparent. Any sample not transparent to electrons looks dark in the TEM which

indicates the necessity for more milling. The cross-section was subsequently analyzed in Philips

119

(FEI) EM420T transmission electron microscope (TEM) with selected area electron diffraction

(SAED).

AFM: The roughness of all the coatings after MOCVD was analyzed by atomic force

microscopy (AFM). The measurements were carried out using Digital Instruments, Dimension

3100, by tapping mode using a silicon tip. The silicon tip with a spring constant of 5 N/m is

attached to the end of a cantilever oscillates at a resonant frequency of 60 kHz. The tip taps the

surface during the bottom of its swing. The root mean square amplitude of the oscillation was

maintained constant by means of a feedback loop. The scan size used for all the samples is 1μm

* 1 μm. The roughness value reported for each coating is an average of RMS roughness values

on four spots chosen randomly on the sample.

XPS: The chemical composition of the coatings was analyzed by X-ray photoelectron

spectroscopy (XPS) with a monochromatic Al Kα source to obtain high-resolution scans for

aluminum 2p, oxygen 1s, and carbon 1s. The binding energies reported in this paper were

obtained by referencing the carbon spectrum to 284.6 eV. It has been reported that acidity of a

surface plays an important role in carbon deposition [20]

. Surface acidity can be characterized by

pyridine adsorption and subsequent examination of the surface by XPS [21,22]

. To analyze the

acidity of the coated surfaces, pyridine adsorption was performed by dipping these coatings in

pure pyridine overnight and evacuating the sample in vacuum at 110-°C for 10 hours

subsequently. XPS on the sample was performed to obtain a high resolution scan for the N 1s. A

deconvolution of the high resolution scan for N1s gives the relative concentration of Lewis and

Brønsted acid sites. The deconvolution of the curve was performed with CasaXPS software. In

the characterization of acidity by pyridine adsorption, the spectrum for N 1s is deconvoluted into

three peaks such that the peak positions correspond to the binding energy values 398.7 eV, 400

120

eV and 401.8 eV. Depending upon the sample pretreatment, if hydroxyl groups are present, these

peaks were assigned to Lewis, weak Brønsted and strong Brønsted sites, respectively. If the

hydroxyl groups are absent, these peaks were assigned to weak Lewis, strong Lewis and

Bronsted acid sites respectively[21,23]

. The presence or absence of hydroxyl group is verified with

the high resolution scan for oxygen 1s.

Temperature-Programmed Oxidation: The substrates containing the solid carbonaceous

deposits are dried under vacuum at 110-°C for 2 hours to remove the adsorbed hydrocarbons.

The amount of solid carbonaceous deposits obtained by thermal stressing was measured by

temperature-programmed oxidation in a RC412 multiphase carbon analyzer [24]

. All the samples

used for analysis are 50 mm long and 3 mm wide. The sample is loaded in a quartz boat and then

heated from 100 °C to 900 °C at ramp rate of 30 °C per minute and held at 900 °C for 5 minutes

under flow of ultra-high purity oxygen. The carbon dioxide produced is measured quantitatively

by an infrared detector. Any CO produced is converted by a copper oxide catalyst to carbon

dioxide prior to the detection. The individual peak positions and peak intensities are

reproducible. The peak temperatures relate to the reactivity / structural order of the deposits

produced by thermal stressing. Hydrogen rich, structurally disordered deposits are oxidized at

lower temperature, while hydrogen lean structurally ordered deposits are oxidized at higher

temperature.


4.4.1 Morphology and Spectroscopic Characterization of Coated Substrates

The surface morphology of the blank substrate AISI304 and its AFM micrograph are

shown, respectively, in Figures 4.4a and 4.4b. The average surface roughness of the blank

121

substrate measured by AFM is 30 nm. The morphology, surface topography and structure of the

alumina coating from aluminum trisecondary butoxide (ATSB) are shown in Figures 4.5a-d. The

morphology of alumina coating produced by ATSB shown in Figure 4.5a suggests that the

coating is porous and has a cauliflower morphology. The evaporator temperature chosen for this

run was 132 °C which is slightly lower than that reported in the literature (138 °C) [8]

. This

temperature was chosen so that the gas phase reactions between precursor molecules are avoided.

It has been suggested that the metal has the tendency to maximize its coordination number which

may lead to the oligomerization of the precursor [25,26]

. But under the conditions chosen, it

appears that oligomerization of the precursor during storage on the shelf and subsequent cracking

of the oligomer in the gas phase increases the precursor concentration in the gas phase.

Therefore, lowering the bubbler temperature did not help reduce the concentration of the

precursor in the gas phase. The higher concentration of the precursor in the gas phase and its

higher rate of decomposition at the reaction temperature of 500 °C led to such a morphology.

The oligomerization of the precursor during storage prevented further experiments to optimize

the conditions needed for a smooth coating. The roughness of the coating measured by atomic

force microscopy in the tapping mode shown in Figure 4.5b is 250 nm. TEM micrograph of the

cross-section of the coating obtained by FIB thinning of the coating in Figure 4.5c shows that the

thickness of the alumina coating ATSB deposited is 600 nm. The bright lines indicated within a

red circle as P in the micrograph Figure 4.5c suggest the presence of pores in the coating

extending from the surface of the film to the film-substrate interface. The bright region indicated

as D in the micrograph Figure 4.5c shows the signs of delamination of the coating from the

substrate. The diffraction pattern of the coating ATSB in Figure 4.5d shows the presence of

diffuse rings suggesting that the coating is amorphous in nature, in agreement with the literature

122

[27]. The surface composition of the coating measured by XPS shows 56 at% of oxygen

corresponding to a binding energy of 531.1 eV and 44 at% of aluminum corresponding to a

binding energy of 74.5 eV. The oxygen to aluminum ratio is 1.27 for the coating which is less

than 1.5 for alumina and 2 for boehmite (AlOOH). In a related study, oxidation of NiAl led to

the formation of an ultrathin alumina film, with oxygen to aluminum ratio of 1.3 [28]

. The authors

explained this non-stoichiometry by suggesting the presence of pyramidal and tetrahedral

coordination on the surface due to coordinative unsaturation. From the O/Al atomic ratio

obtained for ATSB, the surface appears to have coordinative unsaturation. To deduce the

presence of hydroxyl groups in the sample, the oxygen 1s high resolution was deconvoluted into

two segments, the first centered at 531 eV corresponding to the oxide and the second centered at

531.5 eV corresponding to hydroxyl groups. The best fit for the two deconvoluted peaks retraced

the original spectrum indicating the presence of surface hydroxyl groups. Deconvolution of the

high resolution N1s scan for ATSB is shown in Figure 4.5e. The red line is the original spectrum

obtained and the black line is the curve fit for the three deconvoluted peaks. The values of the

area under the nitrogen 1s deconvoluted peak corresponding to the coatings shown in Table 4.2

qualitatively indicates that the surface of the coating obtained from ATSB contains Lewis acid

sites, weak Brønsted acid sites and strong Brønsted acid sites.

Alumina coating (Alacac) was deposited on the surface of AISI304 by the pyrolysis of

another metal organic precursor aluminum acetylacetonate. The SEM micrograph of the Alacac

in Figure 4.6a shows a smooth surface without pores. The roughness of the surface shown by

AFM in Figure 4.6b is 5 nm. The micrograph also shows the presence of spheroidal features

marked as C on the coating in Figure 4.6a which appears to be a characteristic of the coating

obtained from the decomposition of the precursor aluminum acetylacetonate. The same feature is

123

seen as a hump on the surface of the coating in the AFM micrograph. Expansion of a gas

containing a condensable vapor like aluminum secondary butoxide through a subsonic nozzle in

a high pressure system was shown to produce condensate particles whose size distribution

depends on the process parameters like velocity of the fluid, reservoir temperature, pressure and

saturation ratio [29]

. The formation of spherulitic deposits has also been observed during the

deposition of alumina from aluminum acetylacetonate in a previous study and has been attributed

to the condensation of vapors into clusters during the expansion of a precursor-laden gas through

a needle valve [30]

. The system used in this study for the deposition of metal oxide films does not

contain a needle valve. Spherulitic features were not observed in the alumina coating from

aluminum tri-secondary butoxide discussed in the previous section. Therefore, the formation of

spherulitic features during the deposition of alumina from aluminum acetylacetonate appears to

depend upon the decomposition chemistry of the ligand acetylacetonate attached to the metal in

the metal organic precursor providing a carbonaceous environment for the growth of fine

alumina grains that would lead to the formation of a dense pinhole-free alumina film. The

alumina coating obtained from aluminum acetylacetonate appears to aid in the formation of a

smooth and non-porous alumina coating, in comparison to that obtained from aluminum

trisecondary butoxide, making it a better choice as an environmental barrier. The cross-sectional

micrograph of the alumina film Alacac in Figure 4.6c shows the presence of bright regions. It

appears that the density of these regions is less than that of the surroundings due to the relatively

higher transmission of electrons, suggesting the presence of excess hydrogen or carbon in these

regions. The thickness of the coating is 170 nm. Selected area diffraction of the coating shown

in Figure 4.6d suggests that the coating is amorphous. The composition of the alumina coating

Alacac analyzed by X-ray photoelectron spectroscopy after sputtering the surface with argon

124

ions for 30 seconds shows the presence of 56 at% of oxygen, 14 at% of carbon and 31 at% of

aluminum. Examination of the high resolution scan for O 1s shows the presence of surface

hydroxyl groups. A deconvolution of the high resolution scan for carbon 1s spectrum shows the

presence of graphitic carbon with a binding energy of 284.6 eV and carboxyl group with a

binding energy of 288 eV. Temperature programmed oxidation of the alumina coating Alacac

did not produce a strong carbon dioxide signal with a good signal to noise ratio. The presence of

carbon in the coating at these levels has been reported in the literature [9,17,31]

. Deconvolution of

the high resolution scan for N1s for Alacac shown in Figure 4.6e produces only one peak that

traces the acquired spectrum. Hence the black line which was seen in ATSB representing the

curve fit is absent. This suggests that the surface of the alumina coating Alacac predominantly

has weak Brønsted acid sites. Presence of carboxyl groups on the surface of alumina coating

from Alacac appears to increase the concentration of weak Brønsted acid sites of the coating

compared with that obtained from ATSB.

The parameters used for the deposition of zirconia coating on the surface of AISI304 for a

period of 4 h are shown in Table 4.1.The SEM micrograph and the AFM image of zirconia

coating shown in Figure 4.7a and 4.7b suggests that the surface has a granular morphology with

a roughness of 15 nm. Under the conditions mentioned in Table 4.1, the zirconia coating

obtained is more smooth than the alumina coating from ATSB. The cross-section micrograph of

the coating by TEM shown in Figure 4.7c indicates the absence of pores. There are no signs of

delamination in this coating as opposed to that observed in the cross-sectional micrograph of the

ATSB coating at lower magnification as seen in Figure 4.5c. The average thickness of the

coating is 1.3 μm. Analysis of the diffraction pattern of the film deposited at 500 °C shown in

Figure 4.7d, shows the presence of tetragonal phase and reflections corresponding to (002), (202)

125

and (111) crystallographic planes.. This finding is in agreement with the observations reported in

the literature where the films were deposited under the same conditions [32,33]

. A high resolution

scan for zirconium 3d corresponding to a binding energy of 182.2 eV and O 1s corresponding to

a binding energy of 530.9 eV shows that the ratio of O/Zr was 2. The XPS of zirconia coating

after pyridine adsorption did not produce a strong signal to noise ratio for the high resolution N

1s scan due to the weaker adsorption of pyridine on the coating indicating the presence of weak

acidic sites on zirconia and prevented further analysis.

Tantalum oxide is known to be an inert material not capable of any catalytic activity for

dehydrogenation, carbon-oxygen and carbon-sulfur bond cleavage reactions. It was deposited on

AISI304 by the decomposition of tantalum pentaethoxide under the conditions given in Table

4.1. The micrograph of the coating after deposition and the surface roughness of the coating

measured by AFM are shown in Figures 4.8a and 4.8b respectively. The tantalum oxide coating

is smooth with a roughness of 4.5 nm. Tantalum oxide film deposited under the conditions

mentioned in Table 4.1 is known to be amorphous [34]

. High resolution scan of XPS for Ta 4f

centered at 26.7 eV and O 1s centered at 531 eV shows that the O/Ta atomic ratio is 2.5. Due to

the weak adsorption of pyridine on the coating, the signal to noise ratio for N 1s was very weak

preventing further analysis of this coating.

4.4.2 Analysis of Jet-A sample and TPO of the Deposits on Coated and Uncoated

Substrates

Figures 4.9a and 4.9b show the GC-MS and GC-PFPD chromatograms of the Jet-A sample,

respectively. Using the NIST107 library, the GC peaks were identified and labeled as shown in

Figure 4.9a. It shows that the fuel predominantly consists of n-alkanes. By elemental analysis,

126

the concentration of sulfur compounds was found to be 1160 ppm by weight. With comparison to

standards, some of the peaks in the GC-PFPD chromatogram were identified as dimethyl and

trimethyl benzothiophenes as shown in Figure 4.9b. Characterization of sulfur compounds in

aviation fuels by atomic emission detector (GC-AED) [35]

has shown, in the order of increasing

retention times the presence of thiols, sulfides, disulfides (classified as reactive sulfur species)

and methyl-substituted thiophenes and benzothiophenes (classified as non-reactive sulfur

species). The classification of reactivity of the sulfur compounds is based on their tendency to

undergo hydrodesulfurization. Therefore, the unidentified peaks observed in the chromatogram

with shorter retention times may correspond to reactive sulfur species such as sulfides and

disulfides.

Based on the average amount of carbonaceous deposits obtained from triplicate experiments

on the AISI304 and five coated substrates, Figure 4.10 compares the performance of the coatings

to inhibit carbon deposition and metal sulfide formation from Jet-A at 350 °C, 500 psig and

1mL/min for 5 h. The reproducibility data for the TPO profile of solid carbonaceous deposits on

all the substrates are shown in Appendix A. The amount of carbon deposits formed on each

substrate is more than that needed to form a monolayer as seen from the calculations shown in

Appendix B. The amount of solid carbonaceous deposits on alumina coating ATSB obtained

from aluminum trisecondary butoxide is less than half of that formed on the uncoated substrate.

The amount of solid carbonaceous deposits on tantalum oxide is equal to that on the alumina

coating Alacac. Compared with alumina coating ATSB, the alumina coating Alacac is more

effective in reducing the amount of carbonaceous deposits. The amount of solid carbonaceous

deposits formed on zirconia is in between that of ATSB and Alacac. Among the coatings tested,

the amount of solid carbonaceous deposits is the least on the surface of platinum. The nature of

127

each coating with respect to its influence on the effectiveness towards the inhibition of solid

deposit formation is discussed below.

AISI304: The morphology of carbonaceous deposits and the surface degradation of AISI304

due to the formation of metal sulfides with various morphologies are shown in Figure 4.11. The

FESEM micrograph of Figure 4.11a shows the presence of filaments (F), faceted or prismatic

metal sulfides (P) and spherulitic deposits (S) on AISI 304. The x-ray diffraction analysis of the

sample containing these deposits in Figure 4.11b shows the presence of pyrrhotites. Pyrrhotites

are known to have both hexagonal and prismatic crystal habits [36]

. These are seen as hexagonal

crystallites with six- fold symmetry marked as H and prismatic crystallites marked as P in the

micrograph. Fe-Ni-S ternary phase diagram predicts the same [3]

. The TPO profile shown in

Figure 4.11c appears to contain two broad peaks. But upon closer inspection, one may see that

the profile can be deconvoluted to four peaks in the temperature ranges 250 – 400 ºC, 400 – 600

ºC, 550 – 650 ºC and 550 – 720 ºC. These peaks may be assigned in the order of decreasing

reactivity, to hydrogen-rich carbonaceous deposits marked as B in the Figure 4.11a, spherulitic

solid carbon deposits marked as S in Figure 4.11a, small particles of relatively ordered carbons,

and large platelets or films of ordered carbon structures formed by metal catalysis, respectively

[3]. Based on the analysis of the relative amount of deposits on various metal substrates from the

Jet-A sample under the conditions mentioned for thermal stressing, it was suggested that the

formation of carbonaceous deposits depends on the interaction between metals, organosulfur

compounds and the oxygenated intermediates formed in the intermediate regime [3]

.

Alumina ATSB: Figures 4.12a and 4.12b show the SEM micrographs of the carbonaceous

deposits on the coating and the TPO profile of the carbonaceous deposits, respectively. A

comparison of the micrographs in Figure 4.6a and Figure 4.12a shows that the pores of the

128

coatings are filled with carbonaceous deposits. The micrograph Figure 4.12a shows the presence

of bright white particles marked as B and spherulitic deposits marked as S. In the TPO profile,

the evolution of carbon dioxide in the temperature range between 260 °C and 400 °C can be

ascribed to the oxidation of hydrogen rich amorphous carbon seen as bright regions in the

micrograph indicated as B in Figure 4.12a. These deposits may result from the adsorption of the

precursors and their subsequent reactions on the surface. The CO2 peak in the temperature range

between 400 °C and 600 °C can be ascribed to the oxidation of spherulitic deposits that may

nucleate and form in the fluid phase. The presence of pores indicated as red circles in the TEM

micrograph Figure 4.5c suggests that the fuel may access the metal substrate through the pores

normal to the coated surface. Previous studies show the formation of metal sulfide crystallites at

the interface between the alumina coating deposited by the pyrolysis of aluminum trisecondary

butoxide and the substrate, when the coating is exposed to 5% H2S mixture at 450 °C for 24

hours to test its ability to protect the surface against sulphidation [8]

. This suggests that that the

reactants can access the substrate through the pores in the coating. Therefore, the alumina

coating obtained from ATSB under the MOCVD conditions used does not appear to be a good

environmental barrier to protect the substrate surface. The reason for the higher amount of

carbon deposits on the ATSB alumina compared to Alacac alumina may result from a higher

surface area exposed to fuel because of porosity on ATSB alumina coating as further discussed

in the alumina Alacac section below. It is important to note from these observations that the

presence of alumina coating inhibits the formation of catalytic carbon and metal sulfides, which

were observed on the AISI304 surface, if the coating is not compromised.

Alumina Alacac: The morphology of carbonaceous deposits on the coating after thermal

stressing and their TPO profile are shown in Figures 4.13a and 4.13b, respectively. As discussed

129

above, a deconvolution of the profile into two peaks suggests the oxidation of structurally

disordered hydrogen rich carbon deposits in the temperature range between 260 °C and 400 °C

marked as B and spherulitic deposits in the temperature range between 400 °C and 500 °C

marked as S respectively in Figure 4.13a. As shown in Figure 4.10 and the corresponding TPO

profiles in Figure 12b and 13b, the amount of carbonaceous deposits obtained on alumina Alacac

is half of that obtained on alumina ATSB. Considering that the area scanned in the AFM is 1

μm2

during the roughness measurement, the area swept by the tip is 1.5 μm2 for ATSB coating

due to its porosity and 1 μm2 for Alacac coating due to its smooth nature. In other words, the

surface area of the ATSB coating area is 1.5 times that on Alacac coating. It is known that apart

from temperature and pressure, the ratio of surface area of the substrate to the volume of the

deposition space has a significant influence on the chemistry and kinetics of carbon deposition

[37]. The amount of solid carbonaceous deposits on ATSB is two times larger than that on Alacac,

whereas the surface area of alumina ATSB is 1.5 times larger than that on Alacac. This

discrepancy can be attributed to differences in the surface acidity of the two coatings. The

composition of Jet-A was reported to contain predominantly paraffins with lower concentration

of aromatics [38]

. Cracking, dehydrogenation, polymerization and cyclization reactions that could

lead to the formation of carbonaceous solid deposits may proceed on the Brønsted acid sites via

carbocation intermediates under the conditions in the thermal stressing reactor. In addition to the

porous nature of alumina from ATSB, the presence of strong Brønsted acid sites on the surface

of alumina from ATSB relative to Alacac and the coordinative unsaturation of Lewis acid sites

on the surface of ATSB may increase the catalytic activity of the coating to produce more carbon

deposits. Studies on the coking propensity of carbon and alumina supported catalysts during

hydrodesulfurization suggests that surface acidity of carbon is lower than that of alumina [20]

.

130

The presence of weak Brønsted acid sites and carbon on the alumina coating Alacac may be

attributed to its lower activity relative to that on ATSB toward carbon deposition.

Zirconia: From the micrograph in Figure 4.14a, it appears that the coating has not undergone

any degradation during the process of reheating in the thermal stressing reactor after the

deposition process. The TPO profile in Figure 4.14b shows the evolution of carbon dioxide

corresponding to the oxidation of carbonaceous deposits from thermal stressing. Upon closer

observation, the TPO profile can be deconvoluted into two peaks. The first one in the

temperature range between 260 °C and 400 °C can be assigned to the oxidation of structurally

disordered hydrogen rich carbonaceous deposits, appearing as bright regions marked as B in

Figure 4.14a. The second peak in the temperature range between 400 °C and 500 °C can be

attributed to the oxidation of spherulitic deposits marked as S in Figure 4.14a. EDX analysis on

the spherulitic deposits shows the presence of carbon and oxygen.

Tantalum oxide: The micrograph of the carbonaceous deposits on the tantalum oxide coating

surface and the evolution of CO2 profile during TPO is shown in Figures 4.15a and 4.15b,

respectively. The morphology of the coating shown in Figure 4.15a suggests the presence of

hillocks on the tantalum oxide coating. The seam on the coating observed in Figures 4.8a and

4.15a are due to the marks present on the substrate shown in Figure 4.8a formed during the

polishing stage in producing the original foil. As chemical vapor deposition process adopts the

topography of the substrate, the seam is also observed on the coating. The presence of hillocks

shows the signs of compressive stress in the coating. An EDX analysis on these hillocks detected

tantalum and oxygen. The TPO profile of the solid carbonaceous deposits on the coating shows a

broad peak that can be deconvoluted into two small peaks. The oxidation of hydrogen rich

structurally disordered solid carbonaceous deposits seen as bright regions indicated as B in

131

Figure 4.15a gives a CO2 signal in the temperature range between 260 °C and 400 °C. The peak

in the temperature range between 400 °C and 600 °C corresponds to the oxidation of spherulitic

deposits indicated as S in Figure 4.15a. Weaker pyridine adsorption observed on the surface

suggests that the tantalum oxide coating has weaker acidic sites. The TPO results suggest that the

coating is relatively inert and the weaker interaction between oxygenated intermediates and the

coated surface produces fewer deposit forming precursors that eventually led to the formation of

smaller amount of hydrogen rich, structurally disordered deposits and spherulitic deposits.

Platinum coating: Oxygen spillover was postulated in the oxidation of soot with platinum

supported alumina catalysts [39]

. Dissociative adsorption of oxygen on platinum leads to the

formation of monoatomic activated species that oxidizes soot at much lower temperature [40]

. In

the case of carbon deposition from the fuel, dissolved oxygen in the fuel was expected to

spillover on the surface of platinum during thermal stressing and subsequently oxidize the carbon

deposits formed from the fuel. To see the possibility of maintaining a self cleaning surface by

oxidizing the incipient carbonaceous deposits formed from thermal stressing of Jet-A by oxygen

spillover, platinum was deposited on AISI 304 from platinum acetylacetonate under the

conditions given in Table 4.1. It is known that the oxidation of platinum acetylacetonate at 500

°C with UHP oxygen results in the deposition of pure platinum without residual carbon

contamination from the precursor [19,40]

. A high resolution scan for platinum 4f at a binding

energy of 71.9 eV suggests that it is present in metallic form. The morphology of carbonaceous

deposits obtained after thermal stressing of the platinum coating, the TPO profile of

carbonaceous deposits and the EDX elemental map of platinum, carbon and sulfur are shown,

respectively, in Figures 4.16 a-e. The morphology of solid carbonaceous deposits is similar to

that obtained with all other uncoated and coated surfaces. The EDX map suggests that carbon

132

covers patches of the platinum coating whereas sulfur covers the platinum coating completely.

The amount of solid carbonaceous deposits appears to be the least on the platinum coating. EDX

spectrum on various spots of the coating after thermal stressing shown in Figure 4.16 c suggests

the presence of sulfur, carbon, oxygen and platinum on the coating. As the platinum coating is

100 nm thick, it is not surprising to see the signal from the substrate corresponding to iron.

Organosulfur compounds in jet fuel are known to poison platinum surfaces [41]

. A possible

oxygen spillover process during TPO on the surface of platinum was expected to oxidize the

carbonaceous deposits at temperatures less than 300 °C and produce a sharp peak at lower

temperatures. Contrary to this expectation, the first broad peak in the temperature range between

260 °C and 400 °C is attributed to the oxidation of structurally disordered hydrogen rich

carbonaceous deposits marked as B in Figure 4.16a. The second broad peak in the temperature

range between 400 °C and 600 °C is attributed to the oxidation of spherulitic deposits indicated

as S in Figure 4.11a. The peak positions observed on platinum corresponding to the oxidation

temperature in the TPO profile are qualitatively similar to that obtained on other coatings. This

suggests the absence of oxygen spillover due to the poisoning of the platinum coating deposited

on metal surface and hence the lack of catalytic activity of platinum towards the oxidation of

carbon deposits.

4.5 Conclusion

The exposure of AISI 304 to the sulfur compounds and hydrocarbons present in the Jet-

A sample during thermal stressing causes extensive surface degradation because of the formation

of metal sulfides and deposition of carbonaceous solids. Thin films of metal oxides and platinum

deposited by MOCVD on AISI304 substrates block surface reactions that lead to metal sulfide

formation and catalysis of solid carbon deposition during thermal stressing.

133

The effectiveness of the coatings in mitigating carbon deposition decreased in the

following order Platinum > Ta2O5 > alumina from acetyl acetonate > ZrO2 > alumina from

aluminum trisecondary butoxide > AISI304. The amount of solid carbonaceous deposits formed

on the coated surface is less than that on uncoated AISI304 by a factor of 2 for alumina coating

from aluminum trisecondary butoxide, 3.5 for zirconia coating from zirconium acetylacetonate, 4

for alumina coating from aluminum acetylacetonate and tantalum oxide coating from tantalum

pentaethoxide and 5 for platinum coating from platinum acetylacetonate.

The deposition of carbonaceous solids on the coatings can be attributed to the interaction

between oxygenated intermediates formed during thermal stressing and acidic sites on various

metal oxide coatings. The presence of coordinatively unsaturated Lewis acid sites and strong

Brønsted acid sites on the surface of alumina coating from ATSB may explain the higher activity

of this surface relative to other coatings towards the formation of carbonaceous solids. Under the

conditions chosen for the deposition of the oxide films, ATSB formed a porous alumina coating

with pores running from the surface to the substrate. In contrast, acetylacetonate precursors

produced pinhole free alumina (Alacac) and zirconia films. Further, the presence of carbon on

the alumina coating from aluminum acetylacetonate appears to increase the concentration of

weak Brønsted acid sites and reduce the activity of the surface towards carbon deposition.

Platinum coating gave the lowest carbon deposition of all coatings presumably because of

absence of acid sites on the metallic coating. An anticipated benefit from platinum catalysis of

deposit oxidation to maintain a self-cleaning surface was not observed probably because of

sulfur poisoning of Pt surface upon reactions with sulfur compounds in Jet-A.

134

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Aerosol Production. AIChE J. 1988, 34 (4), 551-557.

30. Singh, M. P.; Shivashankar, S. A., Low-pressure MOCVD of Al2O3 films using aluminium

acetylacetonate as precursor: nucleation and growth. Surface and Coatings Technology 2002,

161 (2-3), 135-143.

31. Maruyama, T.; Arai, S., Aluminum-Oxide Thin-Films Prepared by Chemical Vapor-

Deposition from Aluminum Acetylacetonate. Applied Physics Letters 1992, 60 (3), 322-323.

32. Seshu B. Desu; Tian Shi; Kwok, C. K. In Strucutre, Composition, and Properties of MOCVD

ZrO2 Thin Films, Chemical Vapor Deposition of Metals and Ceramics, Materials Research

Society: 1989; pp 349-356.

33. Torres-Huerta, A. M.; Domínguez-Crespo, M. A.; Ramírez-Meneses, E.; Vargas-García, J.

R., MOCVD of zirconium oxide thin films: Synthesis and characterization. Applied Surface

Science 2009, 255 (9), 4792-4795.

34. Kamiyama, S.; Lesaicherre, P.-Y.; Suzuki, H.; Sakai, A.; Nishiyama, I.; Ishitani, A.,

Ultrathin Tantalum Oxide Capacitor Dielectric Layers Fabricated Using Rapid Thermal

Nitridation prior to Low Pressure Chemical Vapor Deposition. Journal of The Electrochemical

Society 1993, 140 (6), 1617-1625.

35. Link, D. D.; Baltrus, J. P.; Rothenberger, K. S., Class- and structure-specific separation,

analysis, and identification techniques for the characterization of the sulfur components of JP-8

aviation fuel. Energy Fuels 2003, 17 (5), 1292-1302.

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Phase and Heterogeneous Surface Reactions. Chemical Vapor Deposition 1998, 4 (4), 151-158.

38. Yu, J. Thermal Decomposition of Hydrocarbons under Near-critical and Supercritical

conditions. Pennsylvania State University-University Park, 1996. Ph. D Thesis.

39. Neri, G.; Bonaccorsi, L.; Donato, A.; Milone, C.; Musolino, M. G.; Visco, A. M., Catalytic

combustion of diesel soot over metal oxide catalysts. Applied Catalysis B: Environmental 1997,

11 (2), 217-231.

40. Kwak, B. S.; First, P. N.; Erbil, A.; Wilkens, B. J.; Budai, J. D.; Chisholm, M. F.; Boatner, L.

A., Study of Epitaxial Platinum Thin-Films Grown by Metalorganic Chemical Vapor-

Deposition. J. Appl. Phys. 1992, 72 (8), 3735-3740.

41. Melchor, A.; Garbowski, E.; Mathieu, M. V.; Primet, M., Sulfur Poisoning of Pt/Al2O3

Catalysts .1. Determination of Sulfur Coverage by Infrared-Spectroscopy. React. Kinet. Catal.

Lett. 1985, 29 (2), 371-377.

138

Table 4.1. Conditions used to deposit metal oxide coatings by MOCVD

Sample ATSB Alacac Zirconia Tantalum oxide Platinum

Reactor Pressure (Torr)

Evaporator Temperature (°C)

Substrate Temperature (°C)

Ar carrier gas flow rate (sccm)

O2 Oxidant flow rate (sccm)

2 6 1 2 2

132 138 160 110 140

500 500 500 500 500

200 200 250 200 100

- - 150 - 10

Table 4.2. Relative areas (%) of the deconvoluted peaks under the N1s high resolution scan

corresponding to the maximum peak positions

Sample Relative areas (%) of the deconvoluted peaks

Binding energy (eV) 398.1 400.6 402.3

ATSB 3 81 16

Alacac 100

139

Figure 4.1. Schematic diagram of the MOCVD set-up used for the deposition of metal oxide

coatings.

Figure 4.2. Flow reactor set-up for thermal stressing experiments with Jet-fuel [3]

.

140

Figure 4.3. Specimen preparation by FIB for TEM examination: SEM micrograph of the coating

after platinum deposition (a), the specimen cross-section after milling to form a wedge (b), and

the specimen fastened to the TEM grid (c).

Figure 4.4. An SEM image (a) and an AFM images (b) of the blank substrate AISI304.

a b

a b

Pt

W 1 2

c

Specime

n

TEM

Cu grid

141

Figure 4.5. An SEM image (a), an AFM image (b), a cross-sectional TEM image (c), a

diffraction pattern (d), and a high resolution scan for N1s after pyridine adsorption on alumina

coating from aluminum trisecondary butoxide (ATSB) (e).

a b

P

c d

e

D

142

Figure 4.6. An SEM images (a), an AFM image (b), a cross-sectional TEM image (c), a

diffraction pattern (d), a high resolution scan for N1s after pyridine adsorption on alumina

coating from aluminum acetylacetonate (Alacac) (e).

a b

d c

e

C

143

Figure 4.7. An SEM image (a), an AFM image (b), a cross-sectional TEM image (c), and a

diffraction pattern (d) of the Zirconia coating from zirconium acetylacetonate.

Figure 4.8. An SEM image (a) and an AFM image (b) of the Tantalum oxide coating from

tantalum pentaethoxide.

a b

a b

c d

144

a)

b)

Figure 4.9. A GC-MS chromatogram of Jet-A with marked peaks for n-alkanes (a), and PFPD

chromatogram with identified sulfur compounds (b).

145

Figure 4.10. Amount of carbon deposits on metal and coated surfaces from Jet-A at 350 °C, 500

psig with a fuel flow rate of 1 mL/min for 5 h.

146

Figure 4.11. An SEM image (a), a X-ray diffractogram (b), and a TPO profile of deposits (c)

formed on AISI304 from Jet-A at 350 °C and 500 psig for 5 h.


350 °C and 500 psig for 5 h on alumina coating from aluminum trisecondary butoxide (ATSB).

S H

P F

a

c

b a S

B

b B

86 μg/cm2

40 μg/cm2

147


350 °C and 500 psig for 5 h on alumina coating from aluminum acetylacetonate (Alacac).


350 °C and 500 psig for 5 h on Zirconia coating.

b

B

S

a

a b

S

B

C

148


350 °C and 500 psig for 5 h on tantalum oxide coating.

a b B

S

S

B

Pt

O

C

S

a b

c

19 μg/cm2

16 μg/cm2

149

Figure 4.16. An SEM image (a), a TPO profile of carbonaceous deposits from Jet-A at 350 °C

and 500 psig for 5 h on platinum coating (b) an EDX spectrum of the sample (c), an EDX map of

platinum (d), an EDX map of carbon (e) and an EDX map of sulfur (f).

d e

f

Platinum map Carbon map

Sulfur map

150

Chapter 5

Characterization of Zirconia Coatings Deposited by MOCVD and Their Effectiveness in

Inhibiting Solid Deposition from Jet Fuel

5.1 Abstract

Zirconia coatings were deposited on AISI 304 by pyrolysis of zirconium acetylacetonate at

two deposition temperatures 450 °C and 500 °C in a low pressure MOCVD reactor to retain the

residual carbon from the incomplete decomposition of the precursor. The coating at 500 °C was

also deposited in the presence of oxygen to completely remove the residual carbon in the coating.

The morphology of the coatings was characterized by scanning electron microscopy and atomic

force microscopy. Temperature-programmed oxidation (TPO), Raman spectra and attenuated

total reflection IR were used for the quantitative and qualitative analysis of the carbon residue

left in the coating. The coating deposited at 450 °C was unstable due to the crack formation

during the process of cooling the reactor from the deposition temperature to room temperature.

Therefore the effectiveness of the coatings deposited at 500 °C in inhibiting solid deposit

formation were tested from thermal stressing of Jet-A containing 1160 ppm by weight of sulfur

at 350 °C and 500 psig for 5 h in a flow reactor with a flow rate of 1 mL/min. By comparing the

amount of carbonaceous deposits from Jet-A on the two coatings deposited at 500 °C in the

presence and absence of oxygen, the carbon present on the surface of the coating does not appear

to affect the activity of the coating towards carbon deposition.

5.2 Introduction

The formation of carbonaceous solid deposits in the fuel handling system of the aircraft

engines due to the exposure of metals to Jet-A in the pyrolytic [1]

and intermediate regime [2]

is of

151

major concern. Surface passivation by oxide coatings inhibits the formation of metal sulfides and

the activity of the metals towards the formation of carbonaceous deposits. The physical

properties like high hardness, low thermal conductivity [3]

, high dielectric constant [4]

, oxygen ion

conductivity, wear resistance and the chemical properties of zirconia, such as corrosion

resistance [5]

have led to many applications of zirconia in thermal barrier coatings, electrolytes

for ceramic fuel cells [6]

, optical coatings [7]

, and fiber reinforced ceramic matrix composites [8]

.

The formation of halides prevents the use of halogenated and fluorinated metal precursors in the

deposition of an environmental barrier coating in aerospace industries. Auto-catalyzed hydrolysis

above 110 °C prevents the use of metal organic precursors such as zirconium isopropoxide and

zirconium tertiary butoxide [8]

. Zirconium acetylacetonate was therefore used as a precursor in

the deposition of zirconia by pyrolysis of the precursor in the MOCVD reactor. The activity of

the carbon surface towards coke deposition is known to vary with the composition of carbon [9]

.

By varying the deposition temperature or using an oxidant, the amount of residual carbon in the

coating and the composition of the coating changes as a function of deposition temperature. The

effectiveness of the residual carbon in coating in inhibiting carbon deposition and metal sulfide

formation from Jet-A at 350 °C and 500 psig was investigated.


5.3.1 MOCVD Experimental Set-up for Zirconia Coating

Figure 5.1 shows the schematic diagram of the experimental set-up used for the deposition of

the zirconia coatings at three deposition temperatures 400 °C, 450 °C and 500 °C in the absence

of oxygen. The coating was also deposited at 500 °C in the presence of UHP oxygen. The

conditions used for deposition were shown, in Table 5.1. AISI 304 substrate purchased from

152

Goodfellow Ltd was cut into foils that are 10 mm long, 3mm wide and 0.9 mm thick, washed in

soap solution, rinsed with acetone and ethanol and dried in vacuum for an hour before the

MOCVD experiment. The foil substrate is inserted inside a stainless tube reactor that is 6.35 mm

(1/4 inch) in diameter and 300 mm long. The edges of the foil are fitted tightly between the walls

of the reactor to maintain a constant substrate temperature. The system is evacuated to a pressure

of less than 0.1 Torr and the reactor tube is then heated to the desired deposition temperature in a

three zone split furnace in the presence of UHP argon flowing through the purge line. The

sublimation temperature in the bubbler was maintained at 160 °C. Once the deposition

temperature and the sublimation temperature reach the respective set point values, the purge line

is shut down. UHP argon flows as a carrier gas flow rate at a flow rate of 250 sccm through the

bubbler. The precursor vapors are transported into the reactor by ultra high purity argon through

the preheating lines which are maintained at a temperature 10 °C higher than the respective

sublimation temperatures. The experiment is conducted for a period of four hours. The coating

at 500 °C was also deposited in the presence of 150 sccm of UHP oxygen to ensure complete

decomposition of the precursor and eliminate the incorporation of residual carbon. At the end of

the experiment, the carrier gas was shut off and the purge argon was flown to cool the reactor.

Once the reactor reaches the ambient temperature, the system is restored to atmospheric pressure.


The coating deposited at 500 °C was tested in the thermal stressing reactor with Jet-A at a

fuel flow rate of 1 mL/min, 350 °C and 500 psig for 5 hours. The details of the thermal stressing

reactor used for isothermal experiments are described elsewhere [10]

. The qualitative analysis of

Jet-A by gas chromatograph-mass spectrometry and pulsed flame photometric detector are

presented in the Figure 4.9a and Figure 4.9b in chapter 4. The stressing experiment is conducted

153

in a 6.35 mm diameter (1/4-in o.d.), 20 cm long glass coated stainless steel reactor. The substrate

is inserted at the bottom of the isothermal glass coated stainless steel reactor. The reactor with

the substrate is heated in the presence of argon to 350 °C at a reactor pressure of 500 psig (3.5

MPa) with the help of a block heater to maintain isothermal conditions along the length of the

reactor and maintained at that temperature for 4 hours to obtain thermal equilibrium. Ultra zero

air is bubbled into the Jet-A reservoir so that it is saturated with dissolved oxygen during the

course of the experiment. The fuel is pumped into the system at 500 psig. It enters the preheating

line of 3.175 mm diameter (1/8-in o.d.) and 2 m in length. The residence time of the fuel in the

preheating line is 6.3 minutes. It is preheated to 260 °C before entering the reactor. The fuel flow

rate, reactor wall temperature and the pressure are maintained at 1 mL/min, 350 °C and 500 psig

for 5 hours. The residence time of the fuel in the reactor is 1.4 minutes. The fuel is maintained in

the fluid phase during the course of experiment. At the end of the experiment, the residual fuel

in the reactor was removed by purging it with argon.

5.3.3 Characterization of Coatings and Carbon Deposits

Morphology The morphology of the coatings after MOCVD, the carbonaceous deposits and

coating after thermal stressing are characterized by field emission scanning electron microscope

JEOL 6700F (FESEM).

Temperature-Programmed Oxidation (TPO) was only performed on the zirconia coating

deposited at 500 °C to measure the amount of carbon in the coating. The coated substrate

containing the solid carbonaceous deposits is dried under vacuum at 110 °C for 2 hours to

remove the adsorbed hydrocarbons. The amount of solid carbonaceous deposits obtained by

thermal stressing was measured by temperature-programmed oxidation in a RC412 multiphase

154

carbon analyzer [11]

. All the samples used for analysis are 50 mm long and 3 mm wide. The

sample is loaded in a quartz boat and then heated from 100 °C to 900 °C at ramp rate of 30 °C

per minute and held at 900 °C for 5 minutes under the flow of ultra high purity oxygen. The

carbon dioxide produced is measured quantitatively by an infrared detector. Any CO produced is

converted by a copper oxide catalyst to carbon dioxide prior to the detection. The individual peak

positions and peak intensities are reproducible. The total amount of solid carbonaceous deposits

obtained on all substrates is reproducible to within 5% of the deposit mass. The peak

temperatures relate to the reactivity / structural order of the deposits produced by thermal

stressing. Hydrogen rich, structurally disordered deposits are oxidized at lower temperature,

while hydrogen lean structurally ordered deposits are oxidized at higher temperature [11]

. The

individual peak positions and peak intensities are reproducible. The total amount of solid

carbonaceous deposits obtained on each substrate is reproducible to within 5% of the deposit

mass.

Raman Spectra Raman shift qualitatively gives an idea about the nature of carbon incorporated

in the coating due to the incomplete decomposition of the precursor and the crystallinity of the

coating. The spectrum of Raman shift was measured for each coating before and after TPO in

WITec confocal Raman spectrometer. The excitation wavelength used to record the spectrum

was 488 nm. The size of the objective aperture was 40x. The time taken to record each spectrum

was 10 seconds. Each spectrum was obtained on three different spots in a sample. The

representative spectrum is shown. The interaction of a photon with a vibrational or rotational

energy level of a molecule results in the gain or loss of its energy. When the energy from the

photon is absorbed by the molecule in the ground state, the energy of the transmitted photon is

less than that of the incident photon resulting in the formation of Stoke’s line. When the energy

155

of the transmitted photon is greater than that of the incident photon due to its interaction with an

excited molecule, it results in the formation of Antistoke’s line. The shift in the frequency is

mainly caused by the vibrational or rotational transition. The spectrum can be acquired only in

substances that experience a net change in the bond polarizability in the presence of an electric

field.

Attenuated Total Reflection Infrared Spectrum Attenuated total reflection method was

chosen to identify the nature of functional groups present on the coating deposited on the AISI

304 with the help of Bruker IFS 66/S FT-IR Spectrometer. Germanium crystal with 80 μm

diameter contact surface was used as a microscopic objective.

X-ray Photoelectron Spectroscopy The composition of the coatings and the nature of carbon on

the coating was found by Kratos Analytical Axis Ultra X-ray photoelectron spectroscopy

instrument. The spot size of the sample was 700 μm * 300 μm. Monochromatic aluminum K

alpha of energy 1486.6 eV was used. Charge correction in the spectrum was made with respect to

aliphatic carbon at 285.0 eV.


5.4.1 Morphology of Zirconia coating

Figure 5.2a-d shows the morphology of the coating and the EDX map of zirconium,

oxygen and carbon respectively deposited at 500 °C. The spherulitic feature on the surface of the

coating is a characteristic of the coating obtained from the acetylacetonate precursor. The EDX

mapping of the coating in the Figure 5.2 b-d shows that the spherulitic feature indicated in red

circle contains zirconium, oxygen and carbon. Expansion of a gas containing a condensable

vapor like aluminum secondary butoxide through a subsonic nozzle from a high pressure region

156

to a low pressure region was shown to produce condensate particles whose size distribution

depends on the process parameters like velocity of the fluid, reservoir temperature, pressure and

saturation ratio [12]

. These particles take the shape of a sphere in order to minimize the surface

free energy. The system used in this study does not have a nozzle or a needle valve. Therefore,

these results suggest that the formation of spherulitic features may be associated with the

decomposition of the acetylacetonate ligand present in zirconium acetylacetonate. The surface of

the coating is smooth, non-porous, and free from cracks.

5.4.2 Raman spectra of Zirconia coatings

Figure 5.3a shows the Raman spectra of the zirconia coating deposited at three different

temperatures. Raman spectra of any carbonaceous solid typically contains first-order regions in

the wavenumber between 1100 cm-1

to 1800 cm-1

and second-order regions between 2500 cm-

1and 3100 cm

-1 [13]. The peak at 1350 cm

-1 suggests the presence of disordered carbonaceous

material in the coating. The broad peak around 2940 cm-1

suggests the presence of C-H

stretching vibrations. The concentration of the C-H is expected to decrease with increase in

deposition temperature from 400 °C to 500 °C due to the more complete decomposition and

fragmentation of the precursor that would result in the incorporation of a lower amount of carbon

in the coating. As expected, the intensity of the peak at 2940 cm-1

decreases with the increasing

deposition temperature. In the first order region, the presence of a peak in 1355 cm-1

suggests

that the carbonaceous material incorporated in the coating due to the incomplete decomposition

of the precursor is disordered. In the second order region, the suppressed peak at 2450 cm-1

and

the broad peaks in the wavenumber range between 2695 cm-1

and 2735 cm-1

support the

implication [14]

. By measuring the intensity of the D peak at 1355 cm-1

and G peak at 1585 cm-1

, it

was observed that the ratio of ID to IG decreases from 0.827 for a deposition temperature of 400

157

°C to 0.796 for a deposition temperature of 450 °C and falls down to 0.773 for a deposition

temperature of 500 °C and falls at. The broad spectra in the wavenumber range between 550 cm-1

and 600 cm-1

correspond to amorphous zirconia [15]

.

Raman spectrum of the zirconia coating after the removal of carbon by temperature-

programmed oxidation is shown in Figure 5.3b. In an inert atmosphere pure zirconia transforms

from the monoclinic phase to the tetragonal phase above 1170 °C. The tetragonal phase

subsequently transforms to the cubic phase at 2370 °C [16]

. Heating the sample in the presence of

oxygen during the temperature programmed oxidation experiment removes the carbon from the

coating and may initiate the phase transformation. The amorphous coating changes its phase to

monoclinic and tetragonal phases. The peaks corresponding to the monoclinic and tetragonal

phases of zirconia are marked as M and T respectively based on the findings in the literature [17]

and shown in Figure 5.3b. The presence of tetragonal phase after heating the sample to 900 °C in

the presence of oxygen is notable.

5.4.3 Infrared spectrum of zirconia coatings

The structure of the precursor zirconium acetylacetonate is shown in Figure 5.4a. The

attenuated total reflection (ATR) IR spectrum of the zirconia coating and the precursor is shown

in Figure 5.4b. The ATR spectrum is qualitatively similar for the zirconia coatings deposited at

three deposition temperatures 400 °C, 450 °C and 500 °C. The wavenumbers of the

corresponding vibration stretches observed for the precursor and coating samples are assigned to

the respective functional groups by comparison to the literature [18]

. The vibration stretch for

hydroxyl groups is typically observed in the wavenumbers ranging between 2900 cm-1

and 3600

cm-1

. As seen in Figure 5.4a, the precursor does not contain any hydroxyl group in its structure

158

and therefore does not show any absorption in the IR spectrum. The presence of carbonyl groups

should show a vibration stretch between 1720 cm-1

and 1760 cm-1

. Even though, these groups

appear to be present in the precursor structure as shown in Figure 5.4a, the spectrum for the

precursor in Figure 5.4b does not show any absorption in this wavelength range. The carbon and

oxygen exist in the precursor as carboxylate group with a resonance stabilized structure where

the electrons in the double bonds are delocalized. This group shows a characteristic absorption at

1535 cm-1

and 1597 cm-1

in the IR spectrum shown in Figure 5.4b. This group also shows

characteristic vibration stretches at 1378 cm-1

and 1430 cm-1

.

The precursor shows absorptions peaks corresponding to the wavenumbers 1060 cm-1

,

1294 cm-1

and 1380 cm-1

. The C-H symmetric bending vibration stretch of the methyl group at

1380 cm-1

overlaps with the vibration stretches corresponding to the carboxylate group at 1378

cm-1

as seen in Figure 5.4b. The other vibration-stretches between 1000 and 1380 cm-1

correspond to the presence of C-C single bond. The band in the wavenumbers between 2860 cm-

1and 2935 cm

-1 present in the coatings suggests the presence of –CH3 and –CH2 groups and

supports the information obtained from Raman spectra. These vibration stretches are present in

both the precursor and the coatings. This suggests that the pyrolytic decomposition of the

precursor in the temperature range between 400 °C and 500 °C produces a coating that contains

carbonaceous materials. Even though the intensity of the IR spectrum corresponding to the

zirconia coating deposited at 500 °C, it is an artifact introduced during the measurement. The

evanescent waves reflected from the sample is collected and analyzed by germanium detector.

The evanescent wave received by the detector is the sum of the intensities reflected from the

film-air interface and film-substrate interface. Ideally, the reflection from the film-air interface

should be zero to identify the functional groups in the sample. Increase in the deposition

159

temperature from 400 °C to 500 °C reduces the carbon content and hydrogen content in the film

due to the more complete decomposition of the precursor and increases the density of the film

due to the higher surface diffusion. Therefore, the refractive index of the film increases with

increasing deposition temperature. The higher refractive index of the film increases the intensity

of IR radiation reflected from the film-air interface and produces a background by overlapping

with that reflected from the film-substrate interface. As scattering increases with increasing

wavenumber, this background also increases for higher wavenumbers. This is predominant in the

spectrum shown in Figure 5.4b for the coating deposited at 500 °C for wavenumbers less than

1400 cm-1

and for wavenumbers between 1800 cm-1

and 2800 cm-1

. The band between 2900 and

3600 cm-1

corresponds to OH stretching. It suggests the presence of different types of hydroxyl

groups. The nature of hydroxyl groups varies with increasing deposition temperature. Studies in

the decomposition of the precursor zirconium acetylacetonate using thermogravimetry indicated

the formation of zirconium hydroxyl acetate at temperatures less than 350 °C during the

pyrolysis process that subsequently decomposes to zirconium oxycarbonate around 450 °C [19]

.

The carbonates have absorption in the wavenumbers 1440 cm-1

and 1530 cm

-1. Due to the

artifacts introduced in the zirconia coating deposited at 500 °C during the measurements, it is

difficult to identify and confirm the presence of such species.

5.4.4 X-ray Photoelectron Spectroscopy of Residual Carbon in the Zirconia coating

The high resolution scan of carbon in the zirconia coating is shown in Figure 5.5. The

binding energies of the deconvoluted peaks are reported after charge correction for the spectrum

with respect to carbon 1s at 284.6 eV. The deconvolution of the high resolution scan for C 1s

suggests the presence of three functional groups. The polymer Poly(acetylacetoxyethyl

methacrylate) is chosen as a reference to identify the various functional groups in the sample.

160

The binding energy maximum corresponding to each component in the sample is compared with

Poly(acetylacetoxyethyl methacrylate) in the XPS reference handbook to identify the functional

groups. The first peak centered at 285 eV suggests the presence of methyl groups (–CH3), the

second centered at 286.6 eV suggests the presence of –CH2 between carboxyl groups. The third

one at 288 eV suggests the presence of carboxyl groups. The nature of functional groups is the

same on all the zirconia coatings deposited at three deposition temperatures. From Figure 5.4a, it

can be seen that the precursor contains these functional groups. It appears that incomplete

decomposition of the precurosor may have resulted in the incorporation of these functional

groups in the coating.

5.4.5 Temperature-Programmed Oxidation

The coatings deposited at 400 °C and 450 °C disintegrated during the process of cooling

the reactor from the deposition temperature to room temperature. Studies show that the

incorporation of non-diamond carbon in the CVD diamond films increases the intrinsic stress in

the coating [18]

. Presence of impurities that do not belong to the host atoms of a material increases

the residual stress of the coating [19]

. Incorporation of carbon due to the incomplete

decomposition of the zirconium acetylacetonate precursor appears to induce significant amount

of intrinsic stress in these zirconia coatings such that there is significant cracking and

disintegration during the cooling process from the deposition temperature to room temperature

after deposition which is clearly seen in Figure 5.6a for the coating deposited at 450 °C. This

prevented further analysis and testing of the coatings in the thermal stressing reactor to evaluate

their performance. Figure 5.7a shows the TPO profile of residual carbon in the zirconia coatings

deposited at 500 °C by the pyrolysis of zirconium acetylacetonate. Deconvolution of the TPO

peak suggests the presence of two types of carbon. The red and the golden yellow curves are

161

deconvoluted peaks shown for visual convenience. The area under each one of them does not

have any physical significance. The graphs suggest that the structurally disordered hydrogen rich

carbonaceous deposit oxidizes in the temperature range between 250 °C and 500 °C. The

structurally more ordered carbonaceous material incorporated in the coating oxidizes in the

temperature range between 400 °C and 600 °C. The coating at 500 °C was, therefore, chosen to

evaluate its performance. Figure 5.8a shows the micrograph of the coating after thermal stressing

with Jet-A. The micrograph shows the presence of spherulitic carbonaceous solid deposits

marked as S and bright regions of structurally disordered hydrogen rich carbonaceous solid

deposits marked as B formed on the coating surface during thermal stressing of Jet-A. The

micrograph also shows the presence of cracks marked as C. These cracks suggest the presence of

residual stress in the coating. It appears that the residual stress is not significant enough to

produce cracks that run perpendicular to the film substrate interface and expose the film substrate

interface. The amount of solid carbonaceous deposits on the coated surface appears to be greater

than that on the uncoated surface at the first sight. The carbon incorporated in the coating during

deposition, due to the incomplete decomposition of the precursor and the carbonaceous solid

deposits from thermal stressing of Jet-A contributes together to the total amount of carbon

dioxide evolved during TPO. Subtracting the former from the total amount shown in Figure 5.8b,

the amount of solid carbonaceous deposits formed during thermal stressing of Jet-A on the

coating is found to be 24 μg per cm2. This is less than that on the uncoated AISI 304 which gives

86 μg per cm2. The micrograph also shows the absence of metal sulfides in the coating. The

amount of carbonaceous deposits on the zirconia coating discussed in Chapter 4 without any

residual carbon is 25 μg per cm2. As the difference falls within the error bar, it suggests that the

162

residual carbon in zirconia coating appears not to affect the activity of the surface towards

carbon deposition from Jet-A.

5.5 Conclusions

Zirconia coating deposited in the temperature range between 400 °C and 500 °C by the

pyrolysis of zirconium acetylacetonate contains significant amount of carbon incorporated in the

coating due to the incomplete decomposition of the precursor as shown by TPO. Characterization

of the coating by Raman spectra, XPS and IR suggests the presence of aliphatic, carbonyl,

carboxyl and hydroxyl groups on the surface of the coatings. The coating deposited at 400 °C

and 450 °C appears to have substantial amount of intrinsic stress possibly due to the

incorporation of carbon in the coating such that the coating disintegrates with significant amount

of surface cracks during the process of cooling from the deposition temperature to the room

temperature. These cracks in the coating would provide access of the fuel to the substrate during

the thermal stressing and therefore are not suitable candidates for an environmental barrier

coating. The coating deposited at 500 °C appears to function better than that deposited at lower

temperatures even though it has also shown signs of cracks formed during the process of

reheating. The amount of solid carbonaceous deposits on the coating is found to be much less

than that on the uncoated surface. The coating deposited at 500 °C also inhibited the formation of

metal sulfides.

163

5.6 References



Research 2008, 47 (23), 9351-9360.

2. Ram Mohan, A.; Eser, S., Analysis of Carbonaceous Solid Deposits from Thermal Oxidative

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165

Figure 5.1. Schematic of the MOCVD setup used for the deposition of coatings

Figure 5.2. (a) SEM of the coating deposited at 500 °C. EDX map of (b) zirconium (c) oxygen

and (d) carbon

a

c

b

d

166

Figure 5.3a. Raman spectra of zirconia coating on AISI 304 before TPO

Figure 5.3b. Raman spectra of zirconia coating on AISI 304 after TPO

Zirconia coating 400



167

Figure 5.4a. Structure of the precursor zirconium acetylacetonate.

Figure 5.4b. Attenuated total reflection infrared spectrum of zirconia coating on AISI 304

168

Figure 5.5. High resolution scan for C 1s in the zirconia coating deposited at 400 °C.

Figure 5.6. SEM micrograph of zirconia coating deposited at 450 °C.

169

Figure 5.7. TPO of residual carbon in the zirconia coatings deposited at 500 °C

Figure 5.8. (a) SEM micrograph of the zirconia coating deposited at 500 °C (b) TPO of the

zirconia coating after thermal stressing with Jet-A at 350 °C, 500 psig and 1 mL/min for 5 hours

carbon representing both carbonaceous deposits from Jet-A and carbon incorporated in the

coating from the precursor.

Zirconia 500

0

Zirconia 400

a

C

S

B

a b

59 μg/cm2

83 μg/cm2

170

Chapter 6

Conclusions, Summary, and Recommendations for Future Work

The principal conclusions and findings of this work are summarized in Sections 6.1 and 6.2

respectively. The recommendations for future work are listed in Section 6.3.

6.1 Conclusions

From the thermal stressing experiments at 350 °C and 500 psig in a flow reactor coupled

with the use of temperature-programmed oxidation to analyze the carbon deposits accumulated

on the foil substrates in the reactor, stainless steel and Inconel alloys were seen to be active

towards carbon deposition from Jet-A. The incipient formation of metal sulfides from the

reaction of sulfur species in the fuel with the Fe and/or Ni on the substrates increases the surface

area available for carbon deposition. This disruption of metal surfaces also exposes the active

metal sites that exhibit catalytic activity towards carbon deposition through dehydrogenation of

adsorbed hydrocarbon species. Metalorganic chemical vapor deposition can be used as a non

line-of-sight deposition technique to coat internal surfaces of tubes of varying diameters. All the

coatings used in this study inhibit the formation of metal sulfides, thereby, eliminate the surface

disruption that increases the area available for carbon deposition and inhibit the formation of

deposits due to the catalytic activity of the base metal. Any smooth, non-porous coating that is

catalytically inactive to dehydrogenation, carbon-oxygen and carbon-sulfur bond cleavage can be

a good environmental barrier by inhibiting the interaction between the fuel and the metal surface.

Organosulfur compounds in jet fuel poison the surface of platinum and prevent the oxygen

spillover process. Therefore, it was not possible to maintain a self-cleaning surface with the

platinum coating.

171

6.2 Summary

This work has focused on the analysis of solid deposits and the effectiveness of

environmental barrier coatings in inhibiting solid deposit formation from Jet-A on seven metal

surfaces in an intermediate regime where both pyrolysis and autoxidation play an important role.

Stressing experiments were conducted in a flow reactor to find out the effect of various metal

surfaces on the formation of solid deposits. A new CVD system was configured to investigate the

possibility of coating tubes of varying diameters used in the fuel handling system. By conducting

experiments with two different precursors and various process conditions for each one of them, a

suitable precursor and the process conditions that would produce a good environmental barrier

coating were identified. Subsequently, four metal oxide coatings and one metal coating were

tested in a thermal stressing reactor to examine their effectiveness in inhibiting solid deposit

formation from Jet-A.

Examination of solid deposits formed from Jet-A on seven metal surfaces with electron

microscopy and temperature-programmed oxidation suggests that the formation of solid

carbonaceous deposits on the metal substrates in the intermediate regime is influenced by

reactive organic sulfides and disulfides in the jet fuel, pyrolysis along with metal catalysis and

the metal sulfide formation. Careful examination of the TPO profile showed the presence of

structurally disordered hydrogen-rich solid carbonaceous deposits, spherulitic deposits, small

particles of relatively ordered carbon and large platelets or films of ordered carbon structures

formed by metal catalysis. X-ray diffraction detected significant amount of pyrrhotites only on

AISI304 and the signal was weak on other metals. A Fe-Ni-S ternary phase diagram was used to

predict the composition of metal sulfides. The phase diagram predicted the formation of

pyrrhotites on iron rich metals and heazlewoodite in nickel rich metal surfaces. Due to the

172

formation of metal sulfides, metals with multiple valence states may be exposed to the

oxygenated intermediates and participate in the decomposition of hydroperoxides through metal-

hydroperoxide complex and other oxidation products formed during liquid phase autoxidation.

Based on the measurement from temperature programmed oxidation of the solid carbonaceous

deposits, the amount of carbon deposition on the alloys increased in the following order AISI316

< AISI 321 ~ AISI 304 < Inconel 600 < AISI 347 < Inconel 718 < FecrAlloy < Inconel 750-X.

The presence of molybdenum, titanium and niobium in smaller amounts in the alloys do not

appear to affect carbon deposition under the experimental conditions. Carbon deposition on

FeCrAlloy, Inconel 600, Inconel 718 and inconel 750-X shows that the formation of metal

sulfides does not necessarily passivate the surface and reduce carbon deposition as might be

expected.

Among the methods available for the deposition of environmental barrier coatings,

metalorganic chemical vapor deposition was identified as a potential process for the deposition

of metal oxide coatings in complex geometries that demand non-line-of-sight deposition. Efforts

to coat the flow passages in injectors for aircraft engines were not successful presumably because

of the presence of tubes with different diameters on the flow path of the fuel in the injectors.

Surface to volume ratio is inversely proportional to the diameter of the tube. In a complex

geometry like a fuel injector where the size of the conduit varies along the tube length, surface to

volume ratio plays an important role in the uniformity of coating thickness. The ratio of

thickness of the coating to the conduit diameter affects the stress normal to the film-substrate

interface. Alumina coatings from aluminum trisecondary butoxide produced at various

conditions have either pores or cracks that made this precursor unsuitable for the deposition of an

environmental barrier coating. The coatings had poor resistance to spallation. Aluminum

173

acetylacetonate is a potential candidate for the deposition of a smooth, pinhole-free and crack-

free environmental barrier coating. XPS indicates the presence of carboxyl groups in the coating.

Platinum deposited from platinum acetylacetonate on the alumina coating (Alacac) has two

oxidation states one corresponding to that of a metal and the other corresponding to that of metal

oxide. The amount of residual carbon in the alumina coating decreases with increasing

deposition temperature from 350 °C – 450 °C.

Thin films of metal oxides and platinum, inhibit the formation of metal sulfides and block

catalytic reactions that form solid carbon deposits. The effectiveness of the coatings in mitigating

carbon deposition decreased in the following order Platinum > Ta2O5 > alumina from aluminum

acetyl acetonate > ZrO2 > alumina from aluminum trisecondary butoxide > AISI304. The

amount of solid carbonaceous deposits formed on the coated surface is less than that on uncoated

AISI304 by a factor of 2 for alumina coating from aluminum trisecondary butoxide, 3.5 for

zirconia coating from zirconium acetylacetonate, 4 for alumina coating from aluminum

acetylacetonate and tantalum oxide coating from tantalum pentaethoxide and 5 for platinum

coating from platinum acetylacetonate.

The amount of deposits on the coatings can be attributed to the interaction between

oxygenated intermediates formed during thermal stressing and the acidic sites on various metal

oxide coatings. The presence of coordinatively unsaturated Lewis acid sites and strong Brønsted

acid sites on the surface of alumina coating from ATSB may explain the higher activity of this

surface relative to other coatings towards formation of more carbonaceous solids. The alumina

coating from aluminum acetylacetonate was found to be more effective in inhibiting deposition

than that from aluminum trisecondary butoxide as the latter is porous and offers more surface

area for carbon deposition. Further, the presence of carbon on the alumina coating from

174

aluminum acetylacetonate appears to increase the concentration of weak Brønsted acid sites and

reduce the activity of the surface towards carbon deposition. Platinum coating gave the lowest

amount of carbon deposits due to the absence of active sites. Sulfur compounds in Jet-A poison

the platinum coating and the catalytic effect of deposit oxidation at lower temperatures to

maintain a self-cleaning surface was not observed.

6.3 Recommendations for Future Work

Thermal cycling of environmental barrier coatings is an important concern which should be

addressed to evaluate the durability and its applicability in fuel handling system of an engine.

The effectiveness of the coating over long duration of time can be tested by depositing the

coating on a thermally grown oxide layer and subjecting the system to thermal cycling. Presence

of carbon on alumina coating from aluminum acetylacetonate appears to change the surface

acidity of the coating. The effect of carbon on alumina coating, from aluminum acetylacetonate,

in changing the surface acidity of the coating can be investigated by changing the variation of

concentration of carbon on the surface. This can be accomplished by depositing the coating by

pyrolysis of aluminum acetylacetonate at various deposition temperatures and examining the

variation in the concentration of Brønsted acid sites by pyridine adsorption. This investigation

will give the variation between the concentration of carbon and the concentration of Brønsted

acid sites. It is observed that the presence of carbon in zirconia coating beyond a certain extent

affects the stability of the coating during the process of reheating to thermal stressing conditions.

Therefore the effect of carbon on the structural stability of the alumina coating should also be

evaluated to find out the effect of Brønsted acid sites on the surface acidity and the amount of

solid carbonaceous deposits formed during thermal stressing.

175

Appendix A

Repeatability Data for TPO Profile of Solid Deposits on Substrate Surfaces


304 from Jet-A at 350 °C and 500 psig for 5 h.



176





177


FeCrAlloy from Jet-A at 350 °C and 500 psig for 5 h.


Inconel 600 from Jet-A at 350 °C and 500 psig for 5 h.

178


Inconel 718 from Jet-A at 350 °C and 500 psig for 5 h.


Inconel 750-X from Jet-A at 350 °C and 500 psig for 5 h.

179

Figure A.9. Reproducibility data for TPO: TPO profile of solid carbonaceous deposits from Jet-

A at 350 °C and 500 psig for 5 h on the alumina coating deposited from aluminum trisecondary

butoxide.


Jet-A at 350 °C and 500 psig for 5 h on the platinum coating deposited from platinum

acetylacetonate on AISI 304.

180


Jet-A at 350 °C and 500 psig for 5 h on the tantalum oxide coating deposited from tantalum

pentaethoxide on AISI 304.


Jet-A at 350 °C and 500 psig for 5 h on the alumina coating deposited from aluminum

acetylacetonate.

181


Jet-A at 350 °C and 500 psig for 5 h on the zirconia coating deposited from zirconium

acetylacetonate.

182

Appendix B

Calculations for the Amount of Carbon Deposits on Substrate Surfaces

Basis: It is assumed that a monolayer of graphite with the (100) crystallographic plane deposited

parallel to the deposition surface whose area is 1 cm2. The distance between carbon atoms in the

(100) crystallographic plane is 0.142 nm.

Area occupied by a carbon atom = (3√3)/4 * (0.142)2 = 0.026194 nm

2.

0.026194 nm2 area is occupied by 1 carbon atom.

1 cm2 area is occupied by 3.8177*10

15 carbon atoms per cm

2 area.

6.023 *1023

atoms of carbon form 1 gmole of carbon.

3.8177*1015

atoms of carbon form 0.63385 * 10-8

gmoles of carbon.

1 gmole of carbon weighs 12 g of carbon.

0.63385 * 10-8

gmoles of carbon weighs 7.6062*10-8

g of carbon.

Therefore, on 1 cm2 deposition surface area, a monolayer weighs 0.076062 μg of carbon. In other

words, 1 μg of carbon when spread as a monolayer occupies 13.14 cm2 surface area.

In this study, the amount of carbon deposits is reported in μg/cm2. The minimum amount of

carbon deposits measured on the substrates is 16 μg/cm2 on the surface of platinum. The

maximum amount of carbon deposits measured is 160 μg/cm2 on the surface of Inconel 750-X. If

it is assumed that the deposited carbon is in the form of graphite, there would be 210 layers

deposited on the surface of platinum and 2105 layers deposited on top of one another on the

surface of Inconel 750-X.

VITA

Arun Ram Mohan

Education

Ph.D.: Energy and Geo-environmental Engineering,

The Pennsylvania State University, University Park, PA

B.Tech.: Chemical Engineering, 2004

National Institute of Technology, Tiruchirappalli, India

Journal Publications

Arun Ram Mohan, Semih Eser, ―Analysis of Carbonaceous Solid Deposits from Thermal

Stressing of Jet-A Fuel on Iron- and Nickel-Based alloy surfaces‖, Ind. Eng. Chem. Res –

ACS 2010, 49, 2722-2730.

Arun Ram Mohan, Semih Eser, ― Effectiveness of Low-Pressure MOCVD Coatings on Metal

Surfaces for the Mitigation of Fouling from Heated Jet Fuel‖ – Accepted for publication in

Ind. Eng. Chem. Res – ACS 2011.

Conference Presentation

Arun Ram Mohan, Semih Eser, ―Inhibition of Carbon Deposition from Jet fuel by surface

coating‖ Fouling Mitigation, ACS and AIChE, New Orleans, USA; 2009.

Research Publications in Proceedings Ramya Venkataraman, Arun Ram Mohan, Eser Semih, ―Alumina coating by chemical vapor

deposition for inhibition of carbonaceous deposits from hydrocarbon decomposition on metal

surfaces.‖ Division of Fuel Chemistry Preprints, American Chemical Society, San Francisco, USA,

August 2006.

Van Nikerk. D, Markley. B, Li. Y, Rodriguez-Santiago. V, Thomson. D, Ram Mohan. A, Elsworth.

D, Jonathan. J.P., Pisupati. S, Song. C, ―Utilization of carbon dioxide from a coal-fired power plant

for the production of value-added products.‖ International Pittsburgh Coal Conference, Pittsburgh,

USA, 2006.

the pennsylvania state university the graduate school

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