On the Degradation Mechanisms of Thermal Barrier Coatings ...

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On the Degradation Mechanisms of Thermal

Barrier Coatings: Effects of bond coat and

substrate

A thesis submitted to the University of Manchester for the degree of

Doctor of Philosophy

in the Faculty of Engineering and Physical Science

2014

LIBERTY TSE SHU WU

School of Materials

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Table of Contents

List of Figures ................................................................................................................ 4

List of Tables ................................................................................................................ 13

Abstract ........................................................................................................................ 15

Declaration ................................................................................................................... 16

Copyright Statement .................................................................................................... 17

Acknowledgement ....................................................................................................... 18

Vita ............................................................................................................................... 19

CHAPTER 1 Literature review .................................................................................... 20

1.1 Background information ........................................................................................ 20

1.1.1 Two types of top coat .......................................................................................... 22

1.1.2 Two types of bond coat ....................................................................................... 23

1.1.3 Oxidation and diffusion behavior in Pt/Pt-Al diffusion bond coats .................... 24

1.1.4 TGO residual stress ............................................................................................. 26

1.1.5 Stress evolution and failure mechanism near TGO interface ............................. 29

1.1.6 The rumpling behavior of Pt modified bond coats .......................................... 34

CHAPTER 2 Interfacial fracture toughness ................................................................ 43

2.1 Background information ........................................................................................ 43

2.2 Overview of the methods for measuring interfacial fracture toughness ................ 44

2.2.1 Conventional methods: Pull-off, double cantilever beam, and scratch adhesion

test, etc. ........................................................................................................................ 46

2.2.2 Indentation .......................................................................................................... 50

2.2.3 Barb ..................................................................................................................... 62

2.2.4 Four point-bending.............................................................................................. 66

2.3 Shortcomings of current interfacial fracture toughness models for indentation .... 70

2.4 Factors affecting interfacial toughness measurement ............................................ 73

CHAPTER 3 Scope and aims of research .................................................................... 75

3.1 Project introduction ................................................................................................ 75

3.2 Considerations and requirements for choosing an adequate adhesion test ............ 76

3.3 Project objectives ................................................................................................... 78

CHAPTER 4 ................................................................................................................ 80

Microstructure parameters affecting interfacial adhesion of Thermal Barrier Coatings

by the EB-PVD method ............................................................................................... 80

4.1 Introduction ............................................................................................................ 80

4.2 Experimental details............................................................................................... 83

4.3 Results .................................................................................................................... 91

4.3.1 Microscopy observation ...................................................................................... 91

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4.3.2 Lifetime and interfacial adhesion of TBC systems ............................................. 95

4.3.3 TGO thicknesses, residual stresses and rumpling at the TGO/bond coat interface

...................................................................................................................................... 97

4.3.4 Phase transformation and Ni/Ti segregation ..................................................... 100

4.4 Discussion ............................................................................................................ 105

4.4.1 Approximation of interfacial fracture toughness Gc ......................................... 105

4.4.2 Linear fitting of crack rate with degradation factors ........................................ 114

4.4.3 Possible degradation mechanisms by ’ formation rate .................................... 119

4.4.4 Stress relaxation by rumpling ........................................................................... 120

4.4.5 The dynamic feature of interfacial fracture toughness...................................... 123

4.5 Conclusion ........................................................................................................... 125

CHAPTER 5 .............................................................................................................. 127

The degradation mechanisms of TBCs: Driving force for spallation versus interfacial

fracture toughness Gc ................................................................................................. 127

5.1 Introduction .......................................................................................................... 127

5.2 Experimental details............................................................................................. 130

5.3 Results .................................................................................................................. 135

5.3.1 Microscopic observation near the TGO/bond coat interface ............................ 135

5.3.2 EBSD mapping analyses near the TGO/bond coat interface ............................ 137

5.3.3 EPMA mapping analysis near TGO/bond coat regions .................................... 144

5.3.4 Hardness by nano-indentation near the TGO/bond coat interface .................... 146

5.3.5 Simulation of phase transformation by Thermo-Calc ....................................... 147

5.4 Discussion ............................................................................................................ 149

5.4.1 Interfacial evolution during thermal exposure .................................................. 149

5.4.2. Phase transformation due to substrate/bond coat chemistry ............................ 152

5.4.3. Phase transformation induced interfacial rumpling ......................................... 154

5.4.4 Driving force for spallation versus interfacial fracture toughness .................... 159

5.5 Conclusion ........................................................................................................... 163

CHAPTER 6 .............................................................................................................. 164

Summary and suggested future work ......................................................................... 164

6.1 Summary .............................................................................................................. 164

6.2 Suggested future work ......................................................................................... 166

References .................................................................................................................. 168

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List of Figures

Fig. 1.1. A schematic diagram showing the multi-layer structure of thermal barrier

coating systems (TBCs) [6]…………………………………………………………..21

Fig. 1.2. A plot showing the improvement in thermal cyclic lifetime of thermal barrier

coating systems (TBCs) with 7 wt% yttria in ZrO2 [7]………………………………22

Fig. 1.3. A comparison of the coating morphology between APS (left) [6] and

EB-PVD (right), showing splat-like and columnar structures, respectively…………23

Fig. 1.4. A comparison of the microstructure between overlay (left) [13] and Pt-Al

diffusion bond coats (right). The diffusion zone is several times thicker in the overlay

bond coat……………………………………………………………………………..24

Fig. 1.5. The evolution of mean frequency shift measured for specimens that failed

after: (a) 190 (b) 225 (c) 207 (d) 237 cycles [28]…………………………………….28

Fig. 1.6. Low stress evolution map showing gradual coalescence of isolated stresses

[28]…………………………………………………………………………………...29

Fig. 1.7. TGO/bond coat interfacial roughness profile of three bond coats: a) LT Pt-Al

b) HT Pt-Al c) Pt diffusion against thermal cycles [35]. The Pt-diffusion bond coat

maintained similar roughness throughout the thermal cycles after 10

cycles…………………………………………………………………………………32

Fig. 1.8. Ternary Ni-Al-Pt phase diagram for three different bond coats: a) LT Pt-Al b)

HT Pt-Al c) Pt- diffusion at 1100°C and 1150°C showing a shift from to ’phase

with decreasing Pt and Al contents [36]……………………………………………...33

Fig. 1.9. The strain response in the bond coat throughout a thermal cycle. Point 7

denotes the onset of isothermal dwelling. The onset of cooling may start at either

point 1 (stress relaxation) or 1’ (residual stress) depending on the duration of

isothermal exposure [39]……………………………………………………………..35

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Fig. 1.10. The stress computed beneath an undulation in the bond coat s a) during

cooling and b) heating with phase transformation; c) during cooling and d) heating

without phase transformation [41]. Bond coats with phase transformation experience a

sharp rise/decline in the plane stresses during cooling and heating………………….38

Fig. 1.11. The difference between 11-22 computed for different bond coat yield

strengths [41]. In general, bond coats with lower yield strengths experience lower

plane stresses during phase transformation due to creep relaxation…………………39

Fig. 1.12. a) The difference between 11-22 computed for different transformation

temperature b) top coat 22 stress computed for different transformation temperature

[40]. Bond coat plane stresses can be relaxed when the transformation temperature is

raised…………………………………………………………………………………41

Fig. 1.13. The average amplitude growth of bond coat surface rumpling as a function

of martensitic transformation temperature showing an increasing rumpling trend with

higher transformation temperature. Note that the Y-axis indicates the % of total

growth that occurred during the transformation [45]………………………………...42

Fig. 2.1. An illustration showing the three pure modes of loading…………………..45

Fig. 2.2. Figures showing the loading angle in barb (left) and pull-off (right)

[48,55]………………………………………………………………………………..46

Fig. 2.3. Illustration of the tensile based direct pull-off test [59]…………………….47

Fig. 2.4. Illustration of the sample and set-up used in the double-cantilever beam

test……………………………………………………………………………………48

Fig. 2.5. a) Pile-up in front of scratching indenter, and b) buckling failure at the

interface of thin coatings [67]………………………………………………………..50

Fig. 2.6. a) the scratching indenter generates a wedge crack some distance ahead; b)

the coating lifted up the wedge crack to cause an interfacial crack c) a through

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thickness crack occurs right in front of the indenter leading to complete spallation

[67]…………………………………………………………………………………...50

Fig. 2.7. Indentation-induced delamination during loading [58]…………………….51

Fig. 2.8. Indentation-induced delamination during unloading [58]………………….51

Fig. 2.9. Figures showing the stages of fracture in a brittle coating with a hard

substrate due to a nano-indenter load [58]…………………………………………...52

Fig. 2.10. Illustration showing the half penny shaped later crack centered under the

indent, which is dependent on film and substrate properties, residual stress, etc

[70]…………………………………………………………………………………...53

Fig. 2.11. SEM image showing a delaminated black diamond film with a

straight-sided buckling behaviour [72]……………………………………………….55

Fig. 2.12. Schematic representation of the relationship between crack length and

diagonal length [79]………………………………………………………………….57

Fig. 2.13. Illustration showing the cross-sectional indentation right on the interface

[51]…………………………………………………………………………………...58

Fig. 2.14. Schematic representation of the crack/indent length – loading relation

[79]…………………………………………………………………………………...60

Fig. 2.15. SEM image showing an example of the crack generated at the interface due

to indentation below [82]. The indentation caused an upward displacement in the

entire coating with the highest point roughly above the central line of the

indent…………………………………………………………………………………61

Fig. 2.16. Illustration of the barb specimen showing a partial removal of the TBC

[55]…………………………………………………………………………………...63

Fig. 2.17. Schematic representation of the barb loading setup (shear mode)

[55]……………...........................................................................................................63

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Fig. 2.18. A plot of the typical loading curve obtained from the barb test on a EB-PVD

TBC system [55]……………………………………………………………………..64

Fig. 2.19. a) Fracture surface at low magnification b) a high-magnification SEM

image showing embedded TGO in the bond coat and the TBC remained on the TGO

[55]…………………………………………………………………………………...65

Fig. 2.20. Illustration of the modified four-point bending specimen and loading setup

with the newly applied stiffener layer for preventing segmentation in the ceramic

coating [54]…………………………………………………………………………..67

Fig. 2.21. A typical force-displacement plot showing the onset of delamination as

indicated by the sudden drop in load P [54]………………………………………….70

Fig. 2.22. Illustration of the two instrinsic failure modes of TBCs

[93]…………………………………………………………………………………...72

Fig. 2.23. Plots showing the dynamic change in the stress, loading angle, and energy

release rate of TBC interfacial degradation [89]……………………………………..73

Fig. 4.1. A comparison of the relative spallation lifetime of commercial TBCs

subjected to thermal cycling, where the X-axis indicates the types of superalloy while

legends indicate bond coats (LCBC = Pt-diffusion, RT22LT = LT Pt-Al, CN91PA =

HT Pt-Al) [26]……………..........................................................................................85

Fig. 4.2. SEM image showing the placement of the Vickers indenter at the TGO/bond

coat interface…………………………………………………………………………88

Fig. 4.3. SEM image showing the crack propagation to the left of the indent……….88

Fig. 4.4. SEM image showing the end of crack propagation to the left of the

indent…………………………………………………………………………………89

Fig. 4.5. A comparison of typical crack lengths between those generated by 9.8 and

4.5N indentation loads, showing differences of roughly two-folds in

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length…………………………………………………………………………………90

Fig. 4.6. Representative interfacial microstructure between the three different bond

coat systems (LT Pt-Al, HT Pt-Al, Pt-diffusion) paired with TMS-82+ superalloy. The

HT Pt-Al bond coat had the least rumpling, yet failed to survive up to

200hrs…………………………...................................................................................93

Fig. 4.7. Representative interfacial microstructure between the three different

superalloys (CMSX-4, TMS-82+, SRR99) paired with LT Pt-Al bond coat (SRR99

failed after 30hrs, and no photos are available for 100 and 200 hrs time trials). The

SRR99 specimen had the most ’phase formation and rumpling, and could not survive

up to 50hrs…………………………............................................................................94

Fig. 4.8. The variation of crack lengths vs. isothermal oxidation (1135°C) hours for the

five different TBC systems (i.e. the prefix TMS-82+ is substrate, while Pt-diffusion is

bond coat) using a load of 4900mN………………………………………………….96

Fig. 4.9. The TGO growth kinetics vs. isothermal oxidation (1135ºC) hours for the five

different TBC systems……………………………......................................................98

Fig. 4.10. The TGO compressive stress vs. isothermal oxidation (1135°C) hours for

the five different TBC systems…………………………………………………….....99

Fig. 4.11. The evolution of interfacial rumpling vs. isothermal oxidation (1135°C)

hours for the five different TBC systems…………………………………………...100

Fig. 4.12. The amount of ’ formation attached to the TGO/bond coat interface vs.

isothermal oxidation (1135°C) hours……………………………………………….101

Fig. 4.13. EDX mapping showing Ti distribution near the TGO/bond coat interface of

the coating after 30hrs of isothermal exposure. An obvious Ti enrichment can be seen

in the bond coat of the SRR99 specimen…………………………………………...102

Fig. 4.14. EPMA mapping showing the elemental mapping of Ni, Ti, Al, and O for

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SRR99 LT Pt-Al after 30hrs of isothermal exposure (the color bars are values in wt

%). The regions of Al and Cr depletion correspond exactly to the enrichment in Ni

and Ti contents…………...........................................................................................103

Fig. 4.15. a) The mole-fraction of different phases existing as a function of Al content

in the SRR99 LT Pt-Al sample obtained from calculation using Thermo-calc, and b)

The comparison of Ti concentration in and ’ phases of the three

alloys………………………………………………………………………………..105

Fig. 4.16. The crack radius (i.e. sum of the length of half indent and crack lengths)

values used in the calculation of Gc in this work…………………………………...109

Fig. 4.17. Micrograph showing the location of the u measurements (6 per

indent)……………………………………………………………………………….111

Fig. 4.18. The upward displacements (i.e. taken as the crack openings at the

TGO/bond coat interface near the corner of indents) used in the calculation of Gc in

this work…………………………………………………………………………….111

Fig. 4.19. The approximated values of Gc for the five TBCs based on the semi-cirular

plate model………………………………………………………………………….112

Fig. 4.20. Micrographs showing the extent of YSZ damage between as-coated TBC

and TBC subjected to 30hrs of isothermal exposure………………………………..113

Fig. 4.21. A quantitative comparison of crack rates amongst the five TBC

specimens……….......................................................................................................116

Fig. 4.22. A quantitative comparison of TGO growth rates amongst the five TBC

specimens…………………………………………………………………………...116

Fig. 4.23. A quantitative comparison of average TGO stresses amongst the five TBC

specimens…………………………………………………………………………...117

Fig. 4.24. A quantitative comparison of TGO rumpling rates amongst the five TBC

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specimens…………………………………………………………………………...117

Fig. 4.25. A quantitative comparison of ’phase formation rates amongst the five TBC

specimens…………………………………………………………………………...118

Fig. 4.26. Plots showing the linear dependency of crack rate on a) ’phase formation

rate, b) TGO thickness growth rate, c) TGO length growth rate (rumpling rate) and, d)

TGO average stress. The crack rate only increases linearly with the rate of ’phase

formation……………………………………………………………………………119

Fig. 4.27. The dependency of stress on the rumpling increase rate (highly linear

dependent)…………………………………………………………………………..122

Fig. 4.28. A semi-quantitative plot illustrating the variation of driving force G (steady

state) and fracture toughness Gc between 30 and 100 hrs of isothermal exposure. △G

gives an indication of the remaining spallation life in the TBC…………………….124

Fig. 5.1. EPMA mapping showing the formation of Ti concentrated ’ in the bond coat

region near the oxide-bond interface of the SRR99 LT Pt-Al (containing traces of

TiO2 in the TGO) and TMS-82+ HT Pt-Al after 30 and 100hrs of isothermal exposure,

respectively………………………………………………………………………….129

Fig. 5.2. SEM micrographs showing a comparison of bond coat regions between 30

and 50hrs isothermally exposed SRR99 LT Pt-Al and CMSX-4 LT Pt-Al specimens,

with a unique ’ formation at the grain boundaries of the SRR99 specimen……….136

Fig. 5.3. SEM micrographs showing a comparison of bond coat regions between 30

and 50 isothermally exposed TMS-82+ HT Pt-Al and TMS-82+ LT Pt-Al specimens,

with a complete martensitic transformation in the HT-Pt-Al bond coat……………137

Fig. 5.4. EBSD mapping showing a comparison of grain morphology of the bond coat

region near the oxide-bond interface of the TMS-82+ HT Pt-Al and TMS-82+ LT

Pt-Al prior to isothermal exposure, with the HT Pt-Al bond coat having much larger

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grains………………………………………………………………………………..138

Fig. 5.5. EBSD mapping showing the lath martensitic structure of the TMS-82+ HT

Pt-Al specimen, in contrast to the phase of the TMS-82+ LT Pt-Al after 30hrs of

isothermal exposure…………………………………………………………………139

Fig. 5.6. EBSD mapping showing the unique ’ precipitation at the grain boundaries

of the SRR LT Pt-Al, in contrast to the clean and uniform phase of the CMSX-4 LT

Pt-Al specimen after 30hrs of isothermal exposure………………………………...139

Fig. 5.7. EBSD mapping showing the degradation in pattern quality of martensitic

region relative to the amount of ’ formation in the bond coat of 100hrs isothermally

exposed CMSX-4 LT Pt-Al…………………………………………………………140

Fig. 5.8. EBSD mapping showing more severe degradation in pattern quality of

martensitic region with increasing amount of ’ formation in the bond coat of 100hrs

isothermally exposed TMS-82+ LT Pt-Al 100hrs…………………………………..141

Fig. 5.9. EBSD mapping showing very severe degradation in pattern quality

associated with even more ’ formation in the bond coat of 100hrs isothermally

exposed TMS-82+ HT Pt-Al………………………………………………………..141

Fig. 5.10. EBSD mapping highlighting the inherent difference between the ’

microstrucuture of the Pt-diffusion bond coat and the microstructure of the LT Pt-Al

bond coat……………………………………………………………………………142

Fig. 5.11. A comparison of the misorientation profile near the oxide-bond coat region

between the TMS-82+ Pt diffusion (top) and SRR99 LT Pt-Al (bottom) after 100hrs

and 30hrs of isothermal exposure, with much sharper peaks in the SRR99 specimen

indicating regions of ’ precipitation………………………………………………..143

Fig. 5.12. EPMA mapping showing the chemistry near the oxide-bond coat regions of

the TMS-82+ Pt-diffusion specimens after 30 and 100hrs of isothermal exposure, with

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uniformly depleted layer of Al and Ti while enriched with Cr……………………..145

Fig. 5.13. EPMA mapping showing a comparison of Cr content near oxide-bond coat

interface between TMS-82+ LT Pt-Al and CMSX-4 LT Pt-Al after 100hrs of

isothermal exposure, with CMSX-4 showing no obvious Ti presence……………..146

Fig. 5.14. A plot showing the hardness values by nano-indentation of various bond

coats near the oxide-bond coat interface, with increasing trend in both hardness values

and standard deviations……………………………………………………………..147

Fig. 5.15. The mole fraction of different phases existing near the TGO/bond coat

interface of the bond coats in a) CMSX-4 LT Pt-Al 100hrs, and b) TMS-82+ LT Pt-Al

50hrs specimens as a function of varying Cr and Ta contents, respectively. These

elements generally help stabilizing phase………………………………………...149

Fig. 5.16. A qualitative illustration of the TGO stresses evolution and interfacial

degradation during isothermal heating and subsequent cooling of the 30hrs

isothermally exposed SRR99LT Pt-Al and the 100hrs isothermally exposed TMS-82+

HT Pt-Al. The SRR99 had more rumpling due to ’ formation at the grain boundary,

while the TMS-82+ had a relatively higher residual stress in the TGO due to

martensitic transformation…………………………..................................................160

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List of Tables

Table 4.1

Nominal composition of the three superalloys used in this work

(wt%)…………………………………………………………………………………84

Table 4.2

Visual representation of the five TBC systems used in this

work……………………………..................................................................................84

Table 4.3

The nominal composition in wt% of the bond coat of the SRR99 specimen as

specified in the calculation of Thermo-Calc relative to its original bulk

composition…………………………………………………………………………104

Table 4.4

The effective in-plane Young’s modulus of the bi-layer coating (YSZ-TGO) after

different isothermal exposures (These values were estimated from Fig. 6 in

Ref.[106])………………...........................................................................................107

Table 4.5

The coating thicknesses (YSZ+TGO thickness) values used for the calculation of Gc

in this work………………………………………………………………………….107

Table 4.6.

A comparison between the Gc reported in Ref. [82] and the TMS-82+ Pt-diffusion

here by approximating u as half-diagonal of the indent mark………………………108

Table 4.7

The comparison between the rankings of performance and contribution factors studied,

showing an exact match between performance and the ’ formation factor………...115

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Table 5.1

Nominal compositions of the three superalloys used in this work (wt%)…………..131

Table 5.2

Visual representation of the five TBC systems used in this work…………………..131

Table 5.3

A comparison between the estimated and actual (EDX scanned) bond coat

compositions in as-coated CMSX-4 specimen……………………………………...134

Table 5.4

The nominal composition in wt% of the CMSX-4 LT Pt-Al and TMS-82+ LT Pt-Al

bond coats as specified in the Thermo-Calc calculation……………………………148

Table 5.5

A summary of the phase evolution, hardness, and TGO stress of all the bond coats

studied between 30 and 100hrs of isothermal exposure…………………………….151

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Abstract

The operating efficiency and reliability of modern jet engines have undergone

significant improvement largely owing to the advances of the materials science over

the past 60 years. The use of both superalloys and TBCs in engine components such

as turbine blades has made it possible for jet engines to operate at higher temperatures,

allowing an optimal balance of fuel economy and thrust power.

Despite the vast improvement in high temperature capability of superalloys, the

utilization of TBCs has brought the concern of coating adhesion during their usage.

TBCs are prone to spallation failure due to interfacial rumpling, which is driven

primarily by thermal coefficient mismatch of the multi-layered structure. Although

interfacial degradation of TBCs has been widely studied by detailed numerical and

analytical models, the predicted results (i.e. stress state and rumpling amplitude) often

deviate from that obtained by experiments. This is largely due to the lack of

consideration of the influence of bond coat and substrate chemistry on the interfacial

evolution of TBC systems. It is only in recent year that more and more study has been

focused on studying the role of chemistry on the interfacial degradation of TBCs.

The purpose of this PhD project is to clarify how the bond coat and substrate chemical

compositions dictate the mechanisms of interfacial degradation, leading to the final

spallation. A cross-sectional indentation technique was utilized to quantitatively

characterize the adhesion of oxide-bond coat interface among 5 systematically

prepared TBC systems. The adhesion of isothermally exposed oxide-bond coat

interface was then correlated with different microstructure parameters, in an attempt

to identify the key parameters controlling the TBC spallation lifetime. EBSD and

EPMA analyses were conducted on the bond coat near the oxide-bond coat interface,

in order to understand the relationship between the key parameters and specific

alloying elements. The results clearly demonstrated that the phase transformation of

bond coat near the oxide-bond coat interface plays the dominant role in the

degradation of interfacial adhesion. Particularly, the co-existence of ’ and martensitic

phases, each having very different thermomechanical response under thermal

exposure, can generate a misfit stress in the TGO layer, and ultimately causes early

TBC spallation. In addition, the phase transformation behavior has been closely

associated with the inherent chemistry of the bond coat and substrate.

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Declaration

No portion of the work referred to in this thesis has been submitted in support of an

application for another degree or qualification of this or any other university or institute

of learning.

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Copyright Statement

(1) The author of this thesis (including any appendices and/or schedules to this thesis)

owns certain copyright or related rights in it (the “Copyright”) and he has given The

University of Manchester certain rights to use such Copyright, including for

administrative purposes.

(2) Copies of this thesis, either in full or in extracts and whether in hard or electronic copy,

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Acknowledgement

First and foremost, I would like to express my sincere gratitude to my doctoral

supervisor, Prof. Ping Xiao, for his immeasurable amount of guidance and support

throughout the course of my PhD study. Without him, this PhD project could not have

been possible in the first place.

I would also like to extend this gratitude to my brother, Dr. Rudder Wu, for his

invaluable contribution to this PhD project. For all these years, he has been my

mentor in academic work and many aspects of life. In addition, I would like to thank

Dr. Xiaofeng Zhao, now working at Shanghai Jiao Tong University, for guiding a

significant portion of this research.

My appreciation also goes to Ms. Hong Gao and Dr. Tadaharu Yokokawa of the

National Institute for Materials Science (NIMS, Japan) for facilitating the analysis of

EBSD and EPMA, respectively. They offered great support on the technical aspects of

the experimental work. I would like to thank the technical staffs in Materials Science

Centre (University of Manchester, UK), especially Mr. Kenneth Gyves and Mr. Andrij

Zadoroshnyj for offering technical support on the cross-sectional indentation work

and raman spectroscopy, respectively. My gratitude extends to Dr. Hiroshi Harada and

Dr. Kyoko Kawagishi of NIMS, Japan, and Mr. Rodney Wing of Chromalloy - UK for

the provision of research facilities and specimens.

Special thanks to my colleagues and friends from the faculty, including Mingwen Bai,

Ying Chen, Kuan-I Lee, Justyna Kulczyk-Malecka, Prince Wang, Geng Xing, Fan

Yang, and Erica Yu. If it wasn’t for their presence, my PhD life would have been a lot

more difficult.

I am especially grateful to my parents for their endless love and support. I thank them

for raising me up and sharing their life philosophy with me.

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Vita

Mr. Liberty Wu was born in Taipei (Taiwan), and spent most of his childhood time there.

At the age of 11, his parents made an important decision to immigrate to Canada. He had

been living in Vancouver (Canada) ever since until receiving a bachelor’s degree in

applied science from the University of British Columbia in 2009. Afterwards, he spent

one year in Tsukuba (Japan), working as a research assistant at the National Institute for

Materials Science (NIMS, Japan). In 2010, he moved onto pursuing a PhD degree in

Materials Science and Engineering at the University of Manchester (UK). He has been

working on the research of thermal barrier coatings under the supervision of Prof. Ping

Xiao. As part of the academic curriculum, he also spent 3 months every year at NIMS

(Japan), carrying out some of his research activities.

-by Prince Wang

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CHAPTER 1

Literature review

1.1 Background information

Ni-based superalloys have been the material choice for the turbine blade in the hot

section of jet engines and land based power generator due to the superior mechanical

properties at high temperatures, particularly the ability to resist creep and fatigue.

Since the invention of jet engine, there has been ongoing effort in raising the turbine

inlet gas temperature, thus, enhancing the fuel efficiency. Inevitably, the alloy

materials used for turbine blades experienced numerous thermal breakdown issues as

the hot gas stream exiting the combustor exceeded the melting temperature of the

metal parts. To address this issue, scientists from NASA first proposed the use of

ceramic coatings on turbine blades in the early 1940, and began adopting this

technology to new prototype planes and rocket applications between 60s and the early

70s [1-3]. From the 1970s to the early 80s, the newly improved plasma sprayed and

EB-PVD thermal barrier coating systems (TBCs) found their utilization in civil jet

engines [4,5]. The TBCs system consists of a yttria-stablized zirconia (YSZ) ceramic

top coating layer deposited onto an underlying substrate, joined with an intermediate

bond coat material, which serves as a reservoir for the preferential formation of an

oxidation resistant alumina thermally grown oxide (TGO) layer (see Fig. 1.1).

21

Fig. 1.1. A schematic diagram showing the multi-layer structure of thermal barrier

coating systems (TBCs) [6].

The top coat, having a tailored composition of approximately 7 wt% yttria, was

determined to maintain the most stabilized tetragonal phase during cyclic thermal

loading, leading to desirable mechanical properties [7,8]. This optimal phase was

reported to have significantly improved TBC lifetime by scientists from NASA in the

70s (see Fig. 1.2). The surface temperature of underlying superalloy of turbines could

decrease by as much as 167°C by using the YSZ coat; thereby, increase the blade

lifetime more than threefold [9].

22

Fig. 1.2. A plot showing the improvement in thermal cyclic lifetime of thermal barrier

coating systems (TBCs) with 7 wt% yttria in ZrO2 [7].

1.1.1 Two types of top coat

The top coat, depending on the types of deposition technique applied, can have

different structural appearances. Usually, the coating layer is deposited using either air

plasma spray method (APS) or electron beam physical vapour deposition (EB-PVD).

The APS coating has a splat-like structure with inter-lamella gaps in-plane with the

substrate; while EB-PVD type gives a characteristic columnar structure with

thru-thickness inter-column gaps (Fig. 1.3). For the APS coating, the lateral strain due

to thermal mismatch between the different layers of coating is accommodated by the

porous structure between the successive layers of splats. The EB-PVD coating, on the

hand, accommodates this lateral strain by its columnar structures with gaps in

between. After exposing to high temperature oxidation, the topcoat-scale interface is

much rougher in the case of APS coating than that of the EB-PVD. The columnar

structure of the EB-PVD coating allows much higher strain compliance between the

different layers; hence, giving a much smoother and adhesive topcoat-scale interface

[9]. Nonetheless, the EB-PVD is known to be more costly while being restricted by its

23

line-of-sight deposition characteristics.

Fig. 1.3. A comparison of the coating morphology between APS (left) [6] and

EB-PVD (right), showing splat-like and columnar structures, respectively.

1.1.2 Two types of bond coat

Similar to the top coat deposition methods, bond coats can also be deposited using

two common techniques: overlay coating or Pt/Pt-Al diffusion bond coats. The former

type generally consists of thermally sprayed metallic alloy in the form of MCrAlX,

where M is usually Ni or a combination of Ni and Co, and the X indicates the minor

element addition, such as silicon, zirconium, hafnium or yttrium. As a result of the

overlaid structure, desirable surface properties could be obtained for specific

applications in coatings of this type. Nonetheless, due to the independent nature of

coating composition, severe inter-diffusion of chemical activities is inevitable

between the bond coat and substrate during thermal cycling (Fig. 1.4).

Unlike the overlay bond coats, the diffusion type utilizes electrodeposition method,

where the bond coat materials are diffused into the superalloy substrate rather than

24

forming a discrete layer. These coating materials have tailored compositions to ensure

a stable formation of alumina scale upon oxidation (Fig. 1.4). For some diffusion

coatings, electrodeposition of thin layers of Pt is done, following a vapour-phase

aluminization. Where as in the so-called “low cost” bond coats patented by

Rolls-Royce and Chromalloy UK [10], the subsequent aluminization is omitted in the

process. It is demonstrated that Pt diffusion process alone could enhance the outward

migration of aluminum, as long as the Pt content is maintained above 15 at % [11].

The major drawback of diffusion coating is that a thicker coating is difficult to

achieve in comparison with the plasma spraying technique of the MCrAlY type [12]

Fig. 1.4. A comparison of the microstructure between overlay (left) [13] and Pt-Al

diffusion bond coats (right). The diffusion zone is several times thicker in the overlay

bond coat.

1.1.3 Oxidation and diffusion behavior in Pt/Pt-Al diffusion bond coats

Due to the porous columnar structure and high ionic diffusivity of YSZ in the ceramic

top coat at high temperature, the ingress of oxygen can readily oxidize the underlying

substrate, leading to catastrophic failure of the turbine structure [14-16]. The

25

oxidation is controlled by the rate of the metal ions migration through the bond coat

or the oxygen diffusion into the sample. The preferred formation of oxide in TBCs is

an adherent and slow growing layer characterized by the parabolic growth, without

internal oxidation [17,18].

During the initial stage of oxidation, a transient form of alumina, θ-Al2O3 is formed

on the surface of Pt/Pt-Al diffusion bond coats. This alumina specie is known to

induce a large compressive stress on the unconstrained sample. Soon, the desirable

-Al2O3 began to transform from the original θ phase, which imposes a tensile stress

against the θ phase. With continuous alumina growth, the overall stress eventually

becomes a steady state of low compression [19].

The interdiffusion of bond coat elements during oxidation will cause a significant

change to the microstructure of bond coat. Interdiffusion zone (IDZ) rich in Ni and

various refractory elements, and voids will appear in the bond coat, which causes a

decrease in load bearing strength of the TBCs [20]. The void formations are thought

to have occurred due to the interdiffusion of Al and outward migration of Ni from

substrate to bond coat. The diffusivities of these two elements differ significantly, and

hence, leading to a flux region of vacancy upon interdiffusion. The vacancy sites then

coalesce to form voids, which is known as Kirkendall effect [21,22].

The primary failure mechanism of turbine components lies in the in-service TBC

spallation due to the coefficient of thermal expansion (CTE) mismatch between the

ceramic top coat of yttria-stabilized zirconia (YSZ) and the bond coat near the

thermally grown oxide (TGO) interface [23,24]. While serving as a reservoir of Al to

26

promote the preferential formation ofα-Al2O3 TGO protective layer, bond coat will

inevitably react with superalloy substrate through outward chemical diffusion from

the substrate at high temperature, which eventually affects the TGO formation

dynamics, microstructure, and interfacial stability [25]. Traditionally, it was thought

that faster TGO growth rate and higher degree of rumpling at the TBC/bond coat

interface imply a shorter spallation lifetime. However, in more recent studies, certain

bond coat/substrate systems showed a relatively longer lifetime despite having either

a thicker TGO and or more rumpling than other TBCs [26]. This phenomenon raises

the incentive to investigate the resisting mechanisms to spallation near the TGO

interface by focusing on the chemical effect of the bond coat and substrate.

Currently, there have been very few systematic attempts at studying the effect of bond

coat and substrate chemistry on TGO spallation behavior. In order to better

understand the interfacial evolution during oxidation, there needs to be a reliable

method of quantifying the interfacial toughness of TBCs subjected to oxidation prior

to complete spallation of the coating. Chapter 2 will briefly review some of the most

commonly used measurement techniques of coating-substrate adhesion, with a

particular focus on the practical aspects of each technique.

1.1.4 TGO residual stress

The evolution of residual stress in the TGO layer of the TBC system has been widely

studied in recent years. One particular method known as photo-luminescence

piezo-spectroscopy (PLPS) has been used to measure the residual stress in the TGO.

Christensen et al [27,28] showed that the stress could be obtained from the

27

stress-induced frequency shift of Cr3+

impurities in solid solution of TGO. This

technique excited much interest due to its non-destructive nature.

Stress is measured by the peak shift relative to an unstrained sapphire reference by

assuming a state of equi-biaxial plane stress, and a random distribution of the

crystallographic texture in the TGO. Based on the relationship obtained by He and

Clarke [29], the frequency shift can be expressed in terms of biaxial stress by the

following equation (Eq. 1-1). Ar laser is used to excite the Cr3+

luminescence spectra,

and their frequency shifts were monitored against cyclic hours (See Fig. 1.5). From

the results of the work by A. Selcuk and A. Atkinson [28], a general trend across

different specimens can be observed, in which the mean compressive stress rises over

the first two thirds of life due to gradual stiffening of the non-planar TGO, and later

decrease toward the end of life by the relaxation of increasingly numerous local

damage events.

∆=5.07 (cm-1GPa-1)(Eq. 1-1)

28

Fig. 1.5. The evolution of mean frequency shift measured for specimens that failed

after: (a) 190 (b) 225 (c) 207 (d) 237 cycles [28].

Low stress mapping indicates isolated stress regions at first, then they gradually

become more numerous and tend to coalesce into larger regions of damage near the

end of life (Fig. 1.6). Early works of PLPS mapping analysis indicate that the

luminescence line-width is much wider than expected, and cannot be explained by

surface roughness alone. Therefore, a constrained deconvolution approach was

developed by several researchers to split up the spectra into two stress contribution; a

high stress range (2.0-4.5GPa) and a low stress range (1- 1.5GPa) [30].

29

Fig. 1.6. Low stress evolution map showing gradual coalescence of isolated stresses

[28].

1.1.5 Stress evolution and failure mechanism near TGO interface

In the PLPS works of Tolpygo and Clarke [31], specimens with smoother surface

(polished) were found to have higher TGO stress (shift) that either maintained

throughout thermal-cycling, or showed no obvious decline compared to those with

rougher bond coat surface. Selcuk et al [30] also indicated that the compressive stress

in a planar TGO due to thermal expansion difference between the alumina and

CMSX-4 substrate is significantly higher than that of a non-planar TGO morphology.

30

Despite the different initial surface morphology, all specimens were found to become

more or less rumpled after thermal-cycling.

Various studies have reported that the TGO stress can fall below expected values by

reasons such as: non-planar TGO morphology; plastic deformation (yield and creep in

the bond coat) during cooling; or damage in the TGO (fracture or detachment from

the bond coat). The failure mechanism of typical Pt-Al bond coats has been widely

studied and various results indicate that the plastic deformation in the bond coat plays

an important role in the type of cracking/stress around the TBC interface [32].

The final fracture path in TBC is believed to be closely related to the surface

roughness of bond coat. Experimental results indicate that on specimens with rough

coating interface, rumpling induced damage near the YSZ interface took the dominant

role in the final spallation. Those with smoother coating interface displayed more

damage near the bond coat/TGO interface. In the latter case, the stress accumulated in

the growing TGO thickness would have contributed as the driving energy for

spallation [33]. A recent study suggests that the failure mechanism is strongly

dependent on the differences in high temperature mechanical properties of the bond

coats [34]. In their experiment, TBCs with Pt-Al bond coats underwent phase

transformation from the as-deposited single phase structure into a two phase (Pt,

Ni)Al and '-(Pt, Ni)3Al during the course of Al depletion in high temperature

oxidation.

It has been proposed that the non-uniform volume change associated with this

transformation could lead to rumpling. However, in the recent study by R.T. Wu et al

31

[35], oxidation result disproves this transformation phenomenon as a major role in

inducing bond coat rumpling. According to their studies, the specimen with LT-Pt-Al

bond coat exhibited the most rumpling, yet its oxidation rate was determined to be the

slowest among the three systems investigated. While on their HT Pt-Al bond coat

specimen with relatively lower Al content, results indicate that its rate of rumpling

was considerably lower than that of the LT Pt-Al system. These findings showed the

opposite of what was expected from phase transformation induced rumpling.

Clearly, the phase transformation accompanied with Al depletion cannot be the sole

mechanism for inducing rumpling. Several studies have pointed to the possible

control of rumpling by the resistance to high temperature plastic deformation of bond

coat structures. In the thermal cycling studies by Wu et al [35], the Pt-diffusion bond

coat, which contained a two phase and 'microstructure similar to that of the base

superalloy, had the least noticeable change in the interfacial waviness against thermal

cycling time (Fig. 1.7). Whereas in the Pt-Al bond coat systems, which underwent a

phase transformation from the as-deposited single phase structure into a two phase

(Pt, Ni)Al and '-(Pt, Ni)3Al, had a very drastic increase in roughness with thermal

cycling [34].

32

Fig. 1.7. TGO/bond coat interfacial roughness profile of three bond coats: a) LT Pt-Al

b) HT Pt-Al c) Pt diffusion against thermal cycles [35]. The Pt-diffusion bond coat

maintained similar roughness throughout the thermal cycles after 10 cycles.

A previous study [35] also suggested that a finely dispersed 'within the matrix can

33

contribute to a high temperature strengthening effect in the Pt-Al systems. The extent

of the 'precipitation is not only dependent on the rate at which Al depletion takes

place by oxidation or interdiffusion, but most importantly how close the composition

is to the /'phase boundary [36]. The HT Pt-Al bond coat, having slightly lower

Pt and Al content near its TGO/bond coat interface after thermal exposure, is

relatively closer to the /'phase boundary than the LT Pt-Al systems (Fig. 1.8). As

a result of a higher volume fraction of precipitated ', HT Pt-Al should have better

creep resistance than that of the LT Pt-Al bond coats at elevated temperature. This

hypothesis is consistent with the experimental results by R.T. Wu et al, in which the

leaner Al containing (HT PT-Al) bond coat exhibited a significantly slower rumpling

rate than the Al-rich LT Pt-Al systems [35]. The authors suggested that the role of

high-temperature mechanical properties in different bond coats, in particular high

temperature plasticity, requires further clarification in order to better understand their

rumpling behavior.

Fig. 1.8. Ternary Ni-Al-Pt phase diagram for three different bond coats: a) LT Pt-Al b)

HT Pt-Al c) Pt- diffusion at 1100°C and 1150°C showing a shift from to ’phase

34

with decreasing Pt and Al contents [36].

1.1.6 The rumpling behavior of Pt modified bond coats

A series of papers by Tolpygo and Clarke was devoted to the understanding of

rumpling mechanisms in Pt modified bond coats [37-39]. Tolpygo et al [39]

explained the dependence of rumpling on the cycle length using a semi-quantitative

plot as shown in Fig. 1.9. Plastic yielding can occur during both cooling and heating

when the bond coat creep strength is low at elevated temperatures (i.e. segments 2-3

an 6-7). Due to an intrinsic asymmetry in the plastic response of the bond coat to

tension and compression, the bond coat is in a state of residual compression at the end

of heating cycle (See point 7). The amount of creep relaxation, therefore, undulation

of the bond coat, will then be dependent on the duration of isothermal dwelling.

Further, Tolpygo et al [39] calculated the creep strain (i.e. surface elongation) of the

bond coat due to creep relaxation of the CTE mismatch from previous thermal cycling

at dwelling temperature, and compared that with the actual surface roughness of the

bond coat observed from experiments. The result suggests that CTE mismatch alone

cannot generate the amount of surface undulation observed. Hence, there must be

other mechanisms of inducing surface undulation, which presumably take place at

elevated temperature ranges of thermal cycling (i.e. isothermal dwelling). This will be

the focus of discussions on rumpling mechanisms in later chapters, as the study only

involves isothermal exposures of TBCs.

35

Fig. 1.9. The strain response in the bond coat throughout a thermal cycle. Point 7

denotes the onset of isothermal dwelling. The onset of cooling may start at either

point 1 (stress relaxation) or 1’ (residual stress) depending on the duration of

isothermal exposure [39].

Several rumpling inducing mechanisms during high temperature dwelling were

proposed by Tolpygo et al [40]. The stress associated with the lateral growth strain of

oxide, which is constrained by the underlying metal substrate, can lead to in-plane

compression of the oxide. By taking into account of all the strains acting on the oxide

layer (i.e. metal-oxide bi-layer interaction) throughout a complete thermal cycle, the

growth strain of the oxide was found to cause rumpling of the underlying bond coat

due to creep relaxation of the strain [40]. The magnitude of the growth strain is

determined by the diffusion rate of the oxide forming ions within the oxide scale [40].

Tolpygo et al [38] also suggested that the bond coat tends to swell during high

36

temperature oxidation, as a result of the intrinsic difference in diffusivities between Al

and Ni (i.e. Kirkendall effect). They found that the wavelength of the undulation

stayed similar throughout thermal cycling, while amplitude continued to grow. This

wavelength is surprisingly similar to the grain size of the bond coat near the surface,

suggesting that rumpling may be induced by an uneven response to swelling between

inner parts of the grain and the regions close to the grain boundary.

Further, the volumetric changes associated with to martensitic and 'phase

transformation during thermal cycling were also investigated for possible role in

inducing rumpling. The 'phase transformation was not a requirement for inducing

rumpling, as bond coat had already form undulation in early thermal cyclic times prior

to any 'phase transformation. However, the authors suggest that 'phase

transformation can promote a further growth in rumpling. For the verification of

martensitic transformation, samples were subjected to thermal cycling in a two-zone

furnace at 1150°C and roughly 750-800°C with holding times of 1hr and 10 minutes,

respectively, for a total of 100 cycles. Since the onset temperature of martensitic

transformation for their specimens was well below 750°C, martensitic transformation

would not occur across those temperature ranges during thermal cycling except for the

final cooling to room temperature. The two-zone thermally exposed samples were

then compared to those that had undergone regular 1hr thermal cycling between

1150°C and room temperature for 100 cycles. The results suggested that both set of

samples achieved the same rumpling magnitude, hence, suggesting that martensitic

transformation has little or no effect on the rumpling accumulation during thermal

cycling [39]. This is, however, contradictory to the modeling work by Glynn et al

described below [41].

37

It has been shown from previous research that with modest quenching from

temperatures above 1000°C, a Ni-rich (61 to 68 at. %) NiAl bond coat can transform

into a tetragonally distorted L10 martensite, instead of the ordered -(Ni,Pt)Al phase

at room temperature. [42,43]. Numerous modeling work by Glynn et al [41,44] was

devoted in understanding the general stress response of the TBC system due to the

phase transformation in the bond coat. Figure 1.10 below depicts the difference in

stress response of the bond coat during cooling and heating with/without the

occurrence of martensitic phase transformation. Fig 1.10a (the case with

transformation) indicated a sharp increase in 11, 33 (in-plane bond coat stress

directly below an undulation) and 22 (normal stress in the bond coat directly below

the undulation) due to the transformation from to martensite during cooling at

around 500°C.

38

Fig. 1.10. The stress computed beneath an undulation in the bond coat s a) during

cooling and b) heating with phase transformation; c) during cooling and d) heating

without phase transformation [41]. Bond coats with phase transformation experience a

sharp rise/decline in the plane stresses during cooling and heating.

The effect of the martensitic transformation on the overall TBC stress is clearly

demonstrated in Figure 1.10, as the stress level can be more than two folds at times

when martensite transformation took place.

39

Glynn et al suggested that overall TBC stress behavior is closely related to the

low-temperature bond coat yielding strength. The stress state in the bond coat directly

below the undulation had been reported to be confined such that y = abs(11-22),

and was plotted against temperature during a typical cyclic loading for four different

yield scenarios (see Figure 1.11). It can be demonstrated from the plot that all bond

coats except the one with the highest yield strength (2000MPa) underwent yielding

during the to martensite transformation upon cooling.

Fig. 1.11. The difference between 11-22 computed for different bond coat yield

strengths [41]. In general, bond coats with lower yield strengths experience lower

plane stresses during phase transformation due to creep relaxation.

40

Among all the scenarios studied, only the 500MPa case underwent a substantial

amount of yielding (see flat part of the curve) upon heating during the reverse

transformation from martensite to at around 620°C. Due to this plastic deformation

of the bond coat, both the 22 tensile stress of bond coat and top coat generated during

previous cooling cycle was much relieved. This was however, not the case for the

1000MPa, as the high tensile stress generated during cooling cannot be relaxed due to

the absence of yielding upon heating.

The importance of phase transformation temperature on the overall TBC stress was

also investigated by Glynn et al using three different scenarios: 600°C, 800°C, and

1000°C by assuming that the to martensite and the reverse transformation to be the

same temperature. As expected, when the transformation took place at around 600°C,

much of the strain was accommodated elastically due to the fact that creep was

minimal at this temperature range; thus, receiving the highest accumulation in bond

coat stress and having the least amount of rumpling (Fig. 1.12 and 1.13).

41

Fig. 1.12. a) The difference between 11-22 computed for different transformation

temperature b) top coat 22 stress computed for different transformation temperature

[40]. Bond coat plane stresses can be relaxed when the transformation temperature is

raised.

It can summed up from the above analysis that a desirable bond coat would have two

attributes: 1) a high/strong bond coat yield strength and 2) a martensite

transformation temperature of no more than 500°C, so that deformation occurs below

the creep range. Having these two attributes could largely alleviate the rumpling

42

growth, consequently, prevent early spallation of TBCs (see Figure 1.13).

Fig. 1.13. The average amplitude growth of bond coat surface rumpling as a function

of martensitic transformation temperature showing an increasing rumpling trend with

higher transformation temperature. Note that the Y-axis indicates the % of total

growth that occurred during the transformation [45].

43

CHAPTER 2

Interfacial fracture toughness

2.1 Background information

For functional coating materials, the structural integrity of the coating layer is of

primary concern during operation. Even though considerable efforts have been

dedicated towards improving the reliability of TBCs, the failure mechanisms are not

fully understood. This is in part, due to the complex nature of the coating structure,

which is further complicated by a failure process that is highly sensitive to processing

technique and pattern of use. During in-service thermal and mechanical loading, a

series of events that include crack nucleation, propagation and coalescence will occur

in most cases between either the top coat and TGO, or the bond coat and TGO, and

sometimes be accompanied by a cohesive failure within one of the layers. The

interfacial stability is thus an overall balance between the crack driving force and the

resistance to crack propagation along the relevant interfaces [26,46,47,48].

It has been proposed that a universal damage parameter on the basis of (apparent)

interfacial fracture toughness is necessary to model the damage accumulation as a

function of loading history. The stress field and mechanical response for an interfacial

crack have been characterized with mathematical solutions. Usually, the interfacial

fracture toughness is being evaluated based on one of the following parameters:

adhesive load, force, stress, intensity factor, or energy release rate [49-51].

Consequently, more than a dozen different testing methods have been developed for

estimating the interfacial fracture toughness of TBC-systems (Demarecaux [52],

44

Vasinota [53], Hofinger [54], Kagawa [55], etc). In these testing techniques, a

controlled delamination of the interfaces is being induced through mechanical loading.

Each approach is more or less bounded to limitations, both experimentally and

material wise.

2.2 Overview of the methods for measuring interfacial fracture

toughness

To critically examine the practicality of various measurement techniques, it is

worthwhile considering the necessary quality for an ideal adhesion test method. As

Rickerby [48] pointed out, the requirement for an ideal test method should bear the

following: (i) non-destructive, (ii) easily adaptable to routine testing of geometrically

complex shapes, (iii) relatively simple to perform and interpret, (iv) amenable to

standardization and automation, (v) reproducible and, if possible, quantitative and (vi)

directly related to coating reliability in specific applications. Back in the time when

this review paper was being written, there were very few if any practical methods that

could fulfill a number of these criteria. Through years of effort, a number of test

approaches have been developed to satisfy some of these requirements, though many

challenges and limitations still exist in using these techniques. Some of the most

common approaches include: (1) traditional methods that rely on direct application of

force to induce coating delamination (2) four point-bending tests of notched

multilayer beams (3) indentation techniques (4) barb test. The main issue with the

current techniques lies in the limitation of loading angle that can be used for testing.

Many structural materials, such as TBC, undergo complicated mixed phase loading

angle at the interface instead of a fixed or pure modes (i.e. mode I opening and mode

II in-plane shear) that are associated with the individual measuring techniques (see

45

Figure 2.1). For instance, the barb and pull-off tests shown in Figure 2.2 are limited to

probing coatings under pure mode II and I loading, respectively. [56,57]. In addition,

other challenges and limitations such as difficulty in sample preparations and

experimental setup also hamper the effectiveness of evaluating the interfacial

mechanical properties.

With the above considerations in mind, the next sections will closely examine the

advantages and disadvantages of some of most commonly used approaches in the

field of interfacial toughness measurement. It is intended, through this review, that

one would become familiar with the applicability of the available techniques, and

develop the key knowledge necessary in adopting them to specific material systems.

Fig. 2.1. An illustration showing the three pure modes of loading.

46

Fig. 2.2. Figures showing the loading angle in barb (left) and pull-off (right) [48,55].

2.2.1 Conventional methods: Pull-off, double cantilever beam, and scratch

adhesion test, etc.

About three decades ago, researchers began to seek and develop methods, which

could be used to measure the interfacial adhesion of modern coating-substrate systems.

Some preliminary methods for determining adhesion have thus been proposed. All

these methods are based on mechanical approach of loading as seen in practical

applications of the pull-off, double cantilever beam, and scratch adhesion test

methods, etc. These tests, though destructive in nature, have set the trend for the

development of interfacial adhesion measurement techniques in the last three decades

[48].

47

Pull-off methods. The term is a generalization of several methods that utilize the

application of a force normal to the coating-substrate interface to quantitatively

determine the practical adhesion. The tape test, being one of the popular methods,

applies a pressure sensitive tape to the coating surface to pull the coating off and

thereafter determines the peel force per unit tape width. Although the test is easy to

carry out, it is only limited to thick and rather weakly adhering coatings due to the

fact that the strength of the bonding tape is limited to about 20MN m-2

. An alternative

to the tape test would be to glue two rods to the coating and substrate ends from

which a pure tensile force can be imposed on the coating to detach it from the

substrate (Fig. 2.3). This method is capable of measuring materials with higher

interfacial fracture toughness; however, alignment must be perfect to ensure uniform

loading across. In practice, pure tensile loading condition is hard to achieve, and

instead, these tests often involve a complex mixture of tensile and shear forces, which

renders the interpretation of the results difficult. Further, there is the possibility of

adhesive glue diffusion through thinner coatings during the test, affecting the

film-substrate interfaces. [58]

Fig. 2.3. Illustration of the tensile based direct pull-off test [59].

48

Double cantilever beam test. Two rectangular arms, one of which is spray coated with the

specimen of interest, are joined together by adhesive epoxy and subjected to

cantilever type loading at the drilled pins using an actuator controlled displacement

(Fig. 2.4).The loading mechanism in the test involves almost pure mode I loading to

the coating bonded between two rigid arms. This unique setup requires sawing a

pre-crack with a crack tip to reaching a depth in excess of the penetration layer of the

coating by adhesive epoxy. Post-test examination of the crack surface is necessary to

confirm whether the specimen failed by adhesive or cohesive, in order to verify if the

experiment is successful or not [60]. This setup works quite well for thick coatings;

however, is not suitable when applied to thin coatings due to the difficulty in

generating the pre-crack and the bonding of the rigid plates may damage the

interfaces to be tested.

Fig. 2.4. Illustration of the sample and set-up used in the double-cantilever beam test.

The energy release rate can be approximated as follow

49

(Eq. 2-1)

where E and t are the Young’s Modulus and thickness of the glass beam; respectively,

and δ is the crack opening, L is the crack length [61].

Scratch adhesion test methods. The scratch adhesion test involves drawing a stylus over the

coating surface with a stepwise or continuously increasing normal force until the

coating delaminates. A critical load, Lc, is defined as the load at which the coating is

removed in a regular way along the whole channel length. The testing result could be

influenced by intrinsic factors (i.e. loading rate, scratching speed, indenter tip radius,

instrument-specific designs, etc.), as well as extrinsic parameters such as substrate and

coating modulus, friction coefficient, surface roughness, etc. As such, those factors,

along with the residual stress in the coating must be accounted when applying the

model used for calculating interfacial fracture toughness. The critical load and the

resulting toughness value could only be used to compare different materials given that

the same failure mode occurs. Buckling failure mode is commonly observed in thin

coatings subjected to the scratch test, whereas in thick coatings, a through thickness

crack is more likely to occur leading to a wedge spallation failure (Fig 2.5,2.6). In the

case of buckling failure, it is difficult to make a reliable comparison between coatings

on different substrates, since significant amount of plastic deformation occurred in the

substrate depending on the material hardness. It is only the case of wedge spallation

that such comparisons between different coatings are meaningful. In general, scratch

test can provide a qualitative assessment of hard coatings on softer substrates;

however, for thin and soft coatings, significant amount of plastic deformation may be

associated with the delamination, making it difficult for an adequate comparison

50

between different materials [62-67].

Fig. 2.5. a) Pile-up in front of scratching indenter, and b) buckling failure at the

interface of thin coatings [67].

Fig. 2.6. a) the scratching indenter generates a wedge crack some distance ahead; b)

the coating lifted up the wedge crack to cause an interfacial crack c) a through

thickness crack occurs right in front of the indenter leading to complete spallation

[67].

2.2.2 Indentation

The majority of the work related to the adhesion assessment by indentation techniques

is based on the notion that the crack initiates at and propagates along the interface

using an indenter loaded from the top. Despite this assumption, there are various types

of models, each of which is associated with a different failure mechanism. Thus, care

51

must be taken to identify the delamination mechanism before choosing an appropriate

model to apply. In general, two types of delamination mechanism can occur during

indentation for soft coatings on hard substrates: (a) compressive stress-induced

delamination during loading, (b) tensile stress triggered delamination during

unloading (Fig. 2.7, 2.8). For specimens with relatively weak interfaces compared to

the coating and substrate, failure tends to occur during loading. In the case where the

interface is relatively stronger, or if the whole coating system tends to dissipate the

induced stress in the form of plastic deformation, then interfacial failure usually

occurs as a single buckling during unloading [58].

Fig. 2.7. Indentation-induced delamination during loading [58].

Fig. 2.8. Indentation-induced delamination during unloading [58].

For indentation of a brittle coating on a hard substrate, the delamination mechanism

can become much more complicated than in the case of soft coatings. Considering

that the coating is more likely to fail by buckling instead of shear delamination, a

slightly modified description of the failure mechanism originally proposed by Li et al

[68] was made in Chen’s work [58]. First, through thickness cracks are generated in

52

the radial direction close to the indenter tip. With further penetration, the crack

openings expand with buckling delamination of the coating to their sides. Lastly,

secondary radial cracks form further away from the indenter, leading to partial or full

spallation of the buckled coating (see Fig. 2.9).

Fig. 2.9. Figures showing the stages of fracture in a brittle coating with a hard

substrate due to a nano-indenter load [58].

The earliest work on the modeling of interfacial toughness of indentation began with

the He et al’s analysis of the crack formation initiating from and out of two dissimilar

elastic solids. The study focuses on determining the crack propagation behavior

through the interface by comparing the relative energy release rate between cracks

that either passes through or gets deflected by the interface. Their work became the

modeling basis for other’s analytical work [69].

Marshall and his co-workers first developed a model that accounts for the energy

53

release rate during delamination by comparing the cases of unbuckled and buckled

coatings. The equation for the strain energy release rate is derived as follows [70]:

(Eq. 2-2)

Where the term α is derived from the slope of the plot of buckling load versus edge

displacement, which is 0.38 for materials with a Poisson’s ration of 1/3. For

non-buckling fracture, α is equivalent to 1, and the residual stress term thus does

not contribute to the energy needed to drive delamination.

(Eq. 2-3)

and the indentation stress is given by

(Eq. 2-4),

where Vi is the indentation volume (which can be estimated from either the

load-displacement curve and indenter geometry or the profile of the indentation

impression) and a is the crack length. The subscript f, i and B denote the film,

indentation, and buckling properties respectively (see Fig. 2.10 below).

Fig. 2.10. Illustration showing the half penny shaped later crack centered under the

54

indent, which is dependent on film and substrate properties, residual stress, etc [70].

Hutchingson and Suo made further development to this model as shown by the

modified equations below. [71]

(Eq. 2-5,2-6)

where σr is the residual stress. This modified model assumes that no strain energy

is dissipated through possible cracks found in the substrate.

In a more recent work, Yeap et al modified this model into two separate equations that

are used to describe both non-buckling and buckling cases [72].

(Eq. 2-7 non-buckling)

(Eq. 2-8 buckling)

where

(Eq. 2-9 critical buckling stress)

55

in which the term Y is a dimensionless constant, as determined to be 1.488 and 1 for

circular buckles and straight-sided buckling respectively. The term t and r are the

coating thickness and half radius of the circular crack front normal to the wedge

indent respectively (Fig. 2.11).

Fig. 2.11. SEM image showing a delaminated black diamond film with a

straight-sided buckling behaviour [72].

The indentation induced stress is given by

(Eq. 2-10)

, and Vi is the volume of plastic indentation as follow

(Eq. 2-11)

56

, while the volume of delaminated material is denoted by the term Vc as follow

(Eq. 2-12)

Ritter J et al developed a new approach for determining the interfacial shear strength

with respect to coating hardness based on contact radius at the onset of delamination

[73].

This method is based on the Matthewson’s stress analysis work of a linearly elastic

coating on a rigid substrate using an axisymmetric rigid indenter [74]. The

experimental apparatus is designed such that a direct in-situ observation of the

debonding process can be recorded along with the measurement of indenter contact

dimension .

Mencik et al estimated the fracture toughness of coating based on the cases of radial

cracking, chipping, and delamination (see Fig. 2.9 before), each having a different

method of calculation. Consequently, they derived the interfacial toughness by

treating the interfacial failure as mode I and cint as the appropriate flaw size at the

interface [75]. From this approach, they obtained an estimated value of 0.18MPa m0.5

for the average interfacial toughness of their hybrid coatings-glass specimens, which

lies within the range of results found by previous work. Despite having a good

accuracy, the difficulty in determining the cint could potentially lead to wide scatter of

results [76,77].

Other miscellaneous indentation techniques

57

In addition to improved models as discussed previously, several special indentation

techniques have also been developed. For instance, it has been proposed to use a

wedge indenter to measure adhesion of thin metal coatings instead of the

axisymmetric indenter such as conical or pyramids type, as the wedge indentation is

relatively easier to induce interfacial failure due to its geometry. This special

technique was then applied to a brittle coating on a ductile substrate using a

stress-analysis approach based on the expanding cavity model [78]. In this model,

radial crack formation is not a problem since there is no formation of tensile hoop

stress in the coating. One disadvantage, however, is that the bending effect in the film

and substrate are completely ignored during the indenting cycle, leading to significant

errors for ductile substrates.

To eliminate the effect from plastic deformation, Choulier [79] and Dal Maschio [80]

proposed using Vickers indentation at the cross section of the coating system close to

the interface, while Chicot et al [51] further modified the technique by introducing a

critical load to crack initiation relationship (Fig. 2.12)

Fig. 2.12. Schematic representation of the relationship between crack length and

58

diagonal length [79].

Based on this linear loading vs. crack relationship, a model was established by

considering the mean geometrical characteristics of the substrate and coating couple,

and that the interface behaves locally like an homogeneous material of which elastic

and plastic properties would result from the contribution of both coating and the

substrate [51].

Mathematical expression of the critical load P

(Eq. 2-13 represents the cracking ability of the interface itself, independent of material

geometry)

The terms in Eq. 1-20 are C=1.854E5, P= load in N, a=crack length generated in m,

and HR and HS=hardness of coating and substrate respectively (see Fig. 2.13). The

exponential n in the above equation is a function of the coating thickness as described

in the linear equation below:

Fig. 2.13. Illustration showing the cross-sectional indentation right on the interface

[51].

59

(Eq. 2-14)

By taking the ln of this relationship as a function of loading values P (N), n can then

be obtained simply as the slope of the plot. Hence, in order to determine the

exponential n, it is necessary to conduct cross-sectional indentation on the same

specimen using several different loads.

Mathematical expression of the apparent interface toughness

(Eq. 2-15)

Where

(Eq. 2-16)

One issue affecting the experimental outcome of the measured interfacial fracture

toughness is the existence of residual stress resulted from the differential contraction

between coating and substrate during cooling. Chicot found that if an annealing

treatment is performed before indentation, the straight lines relating to samples with

different coating thicknesses on bilogarithmic curve of crack length vs. load will

converge onto a restricted area where they all intersect at a point (Fig 2.14). In other

words, there exists a load, independent of the coating thickness that corresponds to the

onset of cracking, also known as cracking ability. This point may be considered as a

characteristic property of the coating-substrate adhesiveness [51].

60

Fig. 2.14. Schematic representation of the crack/indent length – loading relation [79].

Once a crack is formed, the propagation will first go through the zone I, where the

cracks remain in the interface plane. Later moving into zone II as indicated by the

sudden change in the curves of the bilogarithmic diagram (Fig. 2.14). From that point

onward, the crack fronts move into the coating and thus no longer correspond to

delamination of the interface. As shown by Figure 2.12, the crack initiated at a critical

load of 5 N corresponding to a critical crack length of 29µm. Prior to this point, the

curve is simply the impression size of the indenter [51].

Although the cross-sectional indentation method as detailed above is simple, the

difficulty of placing the indent tip accurately on the interface has been known to

hamper the reliability of this technique. This is especially the case for complicated

material systems such as the TBC, where the layers of materials have very different

hardness properties. To address this issue, Wang et al [82] recently came up with a

technique based on a clamped circular plate model, which could measure the

interfacial fracture toughness of TBCs by simply placing the indent in the substrate

close to the interface (Fig. 2.15). The energy release rate G can be obtained with the

61

following equation.

(Eq. 2-17)

where Ec, hc, a, and u are the effective in-plane modulus of the coating, the thickness

of the coating, the crack radius, and the central upward displacement of the coating

due to the indentation (see Fig. 2.15), respectively.

Fig. 2.15. SEM image showing an example of the crack generated at the interface due

to indentation below [82]. The indentation caused an upward displacement in the

entire coating with the highest point roughly above the central line of the indent.

Luminescence mapping, where area of TGO delamination (i.e. stress relaxed region)

could be distinguished, was utilized to account for the crack radius, in addition to

SEM observation. FEA method was utilized to accurately account for the upward

62

displacement term u. The interfacial fracture toughness of TBC was determined to be

approximately 29±9 Jm-2

between 35 and 100 thermal cycles, which is similar to the

results obtained using barb test.

2.2.3 Barb

Kagawa et al [55] has developed a new testing technique to facilitate the evaluation of

interfacial fracture toughness, which is specifically aimed for TBC systems under

predominantly mode II loading. They described a test that obviates many of the

challenges and limitations with the testing techniques previously mentioned (i.e.

limitation in loading angles, sample geometries, coating thickness, etc). The testing

routine, based on barb geometry, is taken from a common approach used to evaluate

the fiber reinforced ceramic and/or metal matrix composites. It has been demonstrated

that the barb test can be useful toward investigating the crack growth behaviour in

mode II loading, leading to coating delamination [55].

The barb test specimen is fabricated from an EB-PVD processed TBC specimen. The

TBC layer is carefully ground away using WC polishing tool, leaving only a portion

of the original ceramic coating at the top. The remaining TBC layers were notched at

distance of 3mm from the end of the specimens for the purpose of defining the length

of crack growth. Two identical specimens prepared this way, were glued together

using adhesive epoxy and alignment tool to form the final testing piece (Fig. 2.16).

63

Fig. 2.16. Illustration of the barb specimen showing a partial removal of the TBC

[55].

Figure 2.17 illustrates the loading fixture equipment used for the barb pullout test.

Sapphire plates were carefully positioned to provide optimal load transfer to the

coating, and to avoid any frictional load transfer to the metallic substrate. A

continuous record of force-displacement response during the loading was kept using a

digital memory scope (ORM 1200, Yokogawa Electric, Tokyo, Japan) [55].

Fig. 2.17. Schematic representation of the barb loading setup (shear mode) [55].

64

From the force-displacement plot, the curve increases linearly until a maximum load,

which corresponds to the first crack initiation as observed by a long-standoff distance

optical microscope. Following the peak load, three to four instantaneous force drops

were typically observed, suggesting that crack propagation occurs incrementally until

the final detachment (Fig. 2.18).

Fig. 2.18. A plot of the typical loading curve obtained from the barb test on a EB-PVD

TBC system [55].

The fracture surface is observed with high-resolution SEM micrographs as shown in

Figure 2.19. Based on the microstructural feature of the fracture surface, the

delamination is shown to proceed essentially at the TGO layer, with the crack

propagation occurring in either the TGO/TBC or the TGO/bond coat interface. A

quantitative analysis using this method is done on the EB-PVD TBC specimen using

the model described by Hutchinson et al [71].

(Eq. 2-18)

65

Fig. 2.19. a) Fracture surface at low magnification b) a high-magnification SEM

image showing embedded TGO in the bond coat and the TBC remained on the TGO

[55].

Where σsub and σtbc are the uniaxial stress in the substrate and coating layer

resepectively. Esub and Etbc are the Young’s modulus of the substrate and the coating,

which are taken as 200GPa and 44Gpa respectively. The stress values are calculated

as follow:

(Eq. 2-19)

Substituting the above into Eq. 2-18 yields.

66

(Eq. 2-20)

An average toughness value of 70J/m2

was obtained using Eq. 2-20. This is slightly

higher than those reported elsewhere by indentation and wedge impression testing

[53,56]. Yukawa et al gave several possible explanations for this phenomena, such as

the effect of residual stress at the interface and the Young’s moduli used in the

analysis, as well as possible friction effect resulting from longer contact zone and

larger normal compressive stress acting on the TBC layer due to the barb testing setup.

Clearly, the effect of TBC properties and delamination mechanisms need to be better

understood in order to yield an accurate analysis in Barb testing. The testing protocol

described by Yukawa et al; however, seems applicable to evaluate the interfacial

toughness of the TBC systems. In a more recent work, Kagawa et al [83] re-analyzed

the barb test measurement based on analytical solutions of a beam theory. It was

determined that the crack actually attains a steady-state at mixed loading angle

between 60-70°, due to a bending moment generated from the support. With this

refined work, an average toughness value of 36J/m2 at approximately 65° loading

angle was obtained, which is more consistent to the value obtained in another

literature work [84].

2.2.4 Four point-bending

Charalambides et al [85] first formulated the critical energy release rate of bi-material

interface through the use of a four-point bending test. This method offers many

advantages such as well-established testing routine and simple sample geometry;

67

however, it is mostly restricted to material composites with relatively high fracture

toughness in the debonding layers due to the difficulty in accounting for vertical

cracking, segmentation. In addition, a critical thickness exists for which the energy

necessary for crack propagation can be stored at the interface [54].

In the work by Hofinger et al [54], a modified version of Charalambides’s method is

presented, which enables one to determine the critical energy release rate for

interfacial delamination of a thin and brittle coating. The issues with segmentation

and critical thickness are avoided by attaching the test specimens with a stiffener. The

specimen of the modified test is shown in Figure 2.20. The specimen is notched in the

center with a symmetric interfacial precrack introduced between the inner loading

span.

Fig. 2.20. Illustration of the modified four-point bending specimen and loading setup

with the newly applied stiffener layer for preventing segmentation in the ceramic

coating [54].

An analytical solution is provided to describe the loading condition as shown by the

setup above.

Mb is the constant bending moment along the cross section.

68

(Eq. 2-21)

where P, b, and l are the applied force, the sample width and half of the difference

between outer and inner span, respectively.

A steady-state energy release rate can be achieved, provided that the crack length is

large compared to the total thickness of coating and stiffener [86]. Base on the Hook’s

law, a modified equation of Charalambides’ steady-state energy release rate can be

formulated as follow:

(Eq. 2-22)

Where the second moments of inertia, I2 is given as

(Eq. 2-23)

and

69

(Eq.2-24)

with

(Eq. 2-25)

Where h, E, and v denote the layer thicknesses, Young’s modulus, and Poisson’s ratio,

respectively, and the subscripts 1, 2, and d refer to the ceramic layer, substrate, and

stiffener. In the analysis, Possison ratio is assumed to be the same for all layers, and

that the stiffener is assumed to be equivalent to the substrate material. A plot of the

normalized energy release rate vs. stiffener to substrate thickness ratio shows that with

increasing thickness ratio hd/h2, the effect of Young’s moduli ratio E2 to E1 on the

energy release rate is decreased accordingly. This implies that by stiffening the top

coat, the energy supplied to the crack growth at the interface is increased under

constant substrate loading conditions.

During the mechanical testing, a force vs. displacement curve was recorded until the

onset of the delamination as indicated by a sudden drop in stiffness from the curve’s

straight line (Fig. 2.21). By knowing the critical load Pc and the corresponding crack

length, the critical energy release rate can be obtained by the analytical equation in Eq.

2-22.

70

Fig. 2.21. A typical force-displacement plot showing the onset of delamination as

indicated by the sudden drop in load P [54].

In the four-point bending work by Thery et al. [87], the interfacial fracture toughness

of an as-coated EB-PVD TBC on a β-NiAl bond coat was measured to be 110 J/m− 2

,

which decreased to a range between 55 and 23 J/m− 2

after 50 and 150 1h thermal

cycles at 1100C°. The as-coated toughness in this case seems to be much lower than

the values reported in the work of the cross-sectional indentation and barb test, while

the value after prolonged thermal cycling is just slightly higher. The possible reasons

for the discrepancy in the as-coated value between the different tests will be briefly

discussed in the following sections.

2.3 Shortcomings of current interfacial fracture toughness models for

indentation

Despite the on-going effort in the development and improvement of new and existing

models, none of the current interfacial fracture toughness models could account for

the complex nature of delamination mechanism, crack morphology, and stress

dissipation. For instance, in all of the stress based models, the mathematical models

could only account for either buckling or non-buckling load alone, without possibly

71

taking into account of mixed modes. In addition, many of these models assume the

crack to be circularly in nature with certain measured radius, where as in most real

situations, crack openings have far more complex shapes, and are therefore difficult to

define.

As a result of the simplicity and assumptions in these models, one tends to either over

or underestimate the toughness values, leading to large discrepancy when comparing

the calculated values to those obtained by others using both similar or different

techniques. Further, there has not been a model yet that could be adequately applied to

a coating of multi-layered structures with non-homogeneous interfacial morphology

such as in the case thermally cycled TBC. So far, most of the existing models have

been developed based on the delamination of a single flat interface between two

dissimilar materials.

One major issue with many of the current interfacial toughness models is that the

amount of alloy creep or plastic deformation is often overlooked [88]. This is

problematic particularly for the indentation on ductile materials. To minimize the

effect of plastic deformation, the stress-analysis model based on Vickers indentation

from the cross-section has been proposed.

Another problem with the current techniques of measuring interfacial fracture

toughness lies in the difficulty in representing the true interfacial toughness values of

interfaces with complicated geometries, such as TBCs. The rumpling of the

TGO/bond coat interface due to thermal cycling makes it difficult to account for the

exact loading phase angle Ψ of the intrinsic TBC failure mode. According to Choi et

72

al [89], the TBC interface undergoes a dynamic loading phase angle Ψ by a

competition between edge and buckle delamination (see Fig. 2.22). The energy

release rate along the straight side and front of the blister crack of a mixed

edge-buckle delamination can be denoted as Gside and Gfront, respectively. The dynamic

energy release rate of intrinsic TBC failure, G, as a result of the mixed mode loading

can be elucidated using Figure 2.23 below by the work of Choi et al [93], where Go

and c are the steady energy release rate parallel to the crack front and the critical

stress for buckling, respectively (see Eq. 2-26,2-27).

Fig. 2.22. Illustration of the two instrinsic failure modes of TBCs [93].

(Eq. 2-26)

(Eq. 2-27)

When o (stress generated at the interface) reaches the c (defined in Eq. 2-27 above),

the G values rises accordingly, especially fast on the side of the blister. Soon after,

Gside reaches the value of Go with a loading angle equivalent of 45°. As o continues

to increase, the side will reach closer to 90 ° (mode II), and the front will be close to

60 °.

73

Fig. 2.23. Plots showing the dynamic change in the stress, loading angle, and energy

release rate of TBC interfacial degradation [89].

2.4 Factors affecting interfacial toughness measurement

As described in the previous section, the interfacial toughness values obtained from

different test methods will obviously be different due to the phase angle applied in the

test setup. For instance, the values obtained using the barb test is slightly higher than

those from the cross-sectional indentation due to the fact that the loading angle in the

barb test is known to be relatively higher (65°- 90°). The calculated fracture toughness

will be higher as a result of the contact and friction associated with shear loading (i.e.

mode II). Although the exact phase is unknown in the cross-sectional indentation test

[82], it is reasonable to assume that the angle is much smaller than the barb test due to

a significant portion of out-of-plane displacements. Hence, the results obtained from

the cross-sectional indentation method can be considered fairly similar to that of the

barb test.

74

The residual stresses in the TGO and YSZ are of great importance when measuring

the interfacial fracture toughness of TBCs. As described in the work by Wang et al

[82], the stored strain energy in the TGO and YSZ would be partially released when

buckling of the TBC occurs, contributing to a greater driving force for coating

spallation. This additional driving force must be accounted by the second term in Eq.

2-28 (see below); otherwise, the calculated interfacial fracture toughness results may

be significantly underestimated.

In addition, the interfacial fracture toughness is likely to be affected by the sintering

and anisotropic microstructure of the YSZ top coat. For instance, in the four point

bending method, inter-columnar fracture and shear displacement were often observed

on specimens subjected to short thermal cycling/isothermal oxidation [87]. These

types of damages, resulted from low sintering level of the YSZ, may significantly

decrease the critical load, Pc, necessary for crack propagation under bending, thus,

yielding a much smaller interfacial fracture toughness value than those obtained from

barb or cross-sectional indentation methods. This issue is largely resolved after

prolonged thermal exposure, as the sintering of the YSZ leads to higher stiffness in

the overall coating.

(Eq.2-28)

Beside the loading phase angle of the interface and residual stresses of coating, other

factors such as sample differences due to coating deposition methods, coating

thicknesses, bond coat structural change, and phase transformation would also give

rise to a difference in the value of interfacial fracture toughness.

75

CHAPTER 3

Scope and aims of research

3.1 Project introduction

There is a wide consensus that the leading degradation mechanism in TBC systems is

associated with the interfacial rumpling or ratcheting of the bond coat during the

course of high temperature oxidation and cooling. This act of distortion resulting from

biaxial compression induces tensile stresses perpendicular to the interface [90-94].

Eventually, crack nucleation and growth can take place at TGO/bond coat and

TGO/Top coat interfaces, as well as within the TGO oxide layer.

Many literatures have identified the TGO growth rate as the key influence in the

lifetime of TBC systems, and that spallation is mainly attributed to the strain energy

lying within the TGO with negligible contribution from the top coat [95,96]. However,

it is evident from recent studies that the bond coat and substrate have clear influence

on the degradation mechanism [97,98]. As pointed out in many previous works,

despite having faster TGO growth rate and more rumpling, certain bond coat and

substrate pairs still outperformed others in lifetime experiments.

To this day, there is still a lack of fundamental understanding of the exact degradation

mechanism leading to rumpling and final spallation. It is only in recent years that

more works have dedicated to the study of the effect of substrate and bond coat

compositions on the lifetime of TBC system. The influence of superalloy composition

76

on TBC lifetime had received very little attention in the past while the life-limiting

factors are yet to be clarified.

3.2 Considerations and requirements for choosing an adequate

adhesion test

Chen et al [58] summarized several key considerations for the selection of adhesion

tests in their review paper on interfacial toughness. To produce accurate results,

plastic deformation in the coating or substrate should be avoided, and the stress

responsible for delamination should be maximized. For stress based analysis, the

geometric parameters, which are raised to high powers, should be accurately

accounted for.

In many instances, an accurate measurement of crack geometry can become tricky

due to the difficulty in defining the crack geometry or propagation path. Also, one

should not overemphasize the contribution of residual stress in overall delamination,

since the driving force depends on the combination of residual stress and the applied

loading [58].

To choose an appropriate testing routine, a number of questions should be taken into

considerations as summarized below:

1. What are the bulk properties of the coating and substrate?

- In most techniques, interfacial delamination relies on the transfer of stress

through the coating or the substrate layers. Hence, these layers should

possess a higher cohesive strength than that of the interface. Bending test

77

is usually not suitable for brittle materials; whereas, the double cantilever

tests could be used due to the lower bending stresses inflicted.

2. How thick is the coating?

- Depending on the coating thickness, there often exists a practical limit at

which certain testing techniques could be used reliably. In pull-off tests,

for example, coatings of less than 10 m cannot be examined since the

defects are too small to reach the critical stress level prior to the failure of

the adhesive glue. For thin coatings, methods that generate buckling

failures such as bulge tests and indentation are the most suitable choices.

On the other hand, thick coatings with high bending stiffness may undergo

through thickness shear cracking upon indentation from the top; however,

the cracks may not divert into interface if the coatings are either brittle or

have high interfacial toughness.

3. How high is the adhesive strength?

- In the case of poor adhesion (Gint < 5Jm-2

), nearly every test could be

utilized to determine the interfacial toughness. As soon as the adhesion

strength increases, the number of testing choices becomes limited,

especially for thin and brittle coatings. To achieve the large stress required

to drive delamination, methods that generate additional stresses through

frictional contact, such as scratch and indentation tests, are the best

choices.

4. Is the residual stress controllable?

78

- Though the stressed overlayer method finds many applications in interfacial

toughness investigation, the addition adhesive layer on top must be

deposited such that the residual stress is large to induce spallation in the

coating of interest, yet the overlayer must have superior adhesion with the

coating than that of the tested interface. This is not always easy to achieve

in many coating/material systems.

5. What type of in-service adhesion failure is involved?

- Ideally, it is best to choose the testing techniques that best mimic the

in-service loading conditions at which the coating fails. For example, the

stressed overlayer method finds useful application in coatings subjected to

frequent thermal cycles and multiple depositions during processing.

3.3 Project objectives

Many recent studies have shown that the rate of TGO growth and interfacial rumpling

are only part of the contributing factors for degradation. There is sufficient evidence

from previous research works suggesting that the alloying contents in the bond coat

and substrate can have significant influence on the interfacial evolution; hence, the

spallation lifetime of TBCs.

In this project, a systematic study of the effect of alloying compositions on high

temperature oxidation behavior is done for TBCs with different industrial Ni-based

superalloys and Pt/Pt-Al diffusion bond coats. The attempt here is to identify the

variation in TGO interfacial evolution between different systems of substrates and

79

bond coats, and determine the key degradation factors in controlling TBC lifetime.

Ultimately, this work aims to clarify and establish the missing link between the

degradation mechanisms and chemistry of the bond coat and substrate.

The proposed work can be summarized by the step-by-step objectives listed below:

1. To measure the interfacial adhesion (semi-quantitative) of the TBCs based

on the cross-sectional indentation technique.

2. To compare the interfacial evolution of the different TBC systems in terms

of microstructure parameters such as, TGO growth kinetic, TGO residual

stress, phase distribution of the bond coat near the TGO/bond coat

interface, etc.

3. To establish a relationship between those microstructure parameters and the

interfacial adhesion of the TBCs, and identify the parameter(s) that has the

most influence on interfacial adhesion.

4. To carry out in-depth analysis of the chemistry and phase transformation of

the bond coat near the TGO/bond coat interface using EDX, EPMA, and

EBSD, etc.

5. To clarify how the key degradation mechanism is associated with the

alloying chemistry of the substrate and bond coat.

80

CHAPTER 4

Microstructure parameters affecting interfacial adhesion of

Thermal Barrier Coatings by the EB-PVD method

4.1 Introduction

A.G. Evans and R.A. Miller et. al [90,99] first suggested that the main intrinsic cause

of TBC spallation is due primarily to thermal coefficient mismatch between the

ceramic-metal layers and bond coat oxidation. They also paved the way toward the

development of TBC life models by expressing the critical failure criterion as a

function of the strain energy accumulation near the TBC interface. There is now a

wide consensus that the leading degradation mechanism in thermal barrier coating

systems (TBCs) is associated with the interfacial rumpling of the bond coat during the

course of high temperature oxidation and subsequent cooling, which arises mainly

from mismatch in the coefficient of thermal expansion between the multilayered

structures of the TBCs [23,24]. This act of distortion, resulting from biaxial

compression of the thermally grown oxide (TGO) layer, induces tensile stresses

perpendicular to the interface and eventually leads to spallation [90,91-94].

EB-PVD has been the choice for depositing the state-of-the art ceramic top coat for

many years; while thermal sprayed (i.e. air plasma spraying) coating methods are still

widely used in non-moving components of the jet engine and land-based turbines. It

should be noted that the phenomena observed and mechanisms proposed in this work

are specifically for TBC systems with EB-PVD top coat and diffusional bond coat,

81

and thus, may not be fully relevant to coatings made by thermal spraying and other

methods due to fundamental differences in coating morphology and chemistry.

As more industrial grade TBCs are being developed nowadays, the focus of the latest

TBC research has been shifted toward the comparison of different coating systems. It

is only in recent years that the influence of bond coats and substrates on the modes of

degradation is being studied and compared between different TBCs, in an effort to

understand the key factors affecting the TBC spallation lifetime [26,35,100-102].

Despite the significant progress made in the study of TGO residual stresses and

modelling of TBC failure little work has taken into account of how these stresses and

failure modes could be influenced by the combination of bond coat and substrate alloy.

While the modelling community [23,24,34,103] made detailed numerical description

of the interfacial undulation and corresponding stress state during cyclic oxidation,

their models tend to overlook the dependence of these values on the complicated

nature of bond coat chemical diffusion/phase transformation, etc. Hence, their

prediction of TBC lifetime expressed in terms of a failure criterion (i.e. TGO

thickness) could deviate largely from that obtained in real experiment. So far, there

has not been a successful attempt to come up with a model that could accurately

predict the lifetime of a state-of-the art TBC.

In order to gain a better understanding of the influence of substrate and bond coat

compositions on the degradation mechanisms of TBCs, it is necessary to adopt one of

the aforementioned techniques in Chapter 2 for measuring the TGO/bond coat

interfacial fracture toughness prior to spallation. Conventional methods (i.e. pull-off,

double cantilever beam, etc.) cannot be used to quantify the interfacial fracture

82

toughness due to limitations such as coating thickness and inherent coating adhesive

strength of TBCs. There have been few successful measurements of the interfacial

fracture toughness of TBCs so far, using techniques such as four-point bending,

cross-sectional indentation, and barb test, etc. However, as discussed in section 2.2.3,

2.2.4 and much of the last chapter, these techniques either require tedious

experimental setup (i.e. four-point bending) or extensive geometric/material property

measurement routines.

Considering that a large number of TBC systems will be studied in this project, it is

necessary to adopt a simple yet effective approach to measuring the interfacial

toughness of TBCs. Of all the techniques reviewed in Chapter 2, the cross-sectional

indentation right on the interface developed by Choulier and Dal Maschio [79,80]

seems to satisfy this demand; hence, it will be utilized in this project to provide a

semi-quantitative measurement of interfacial adhesion of TBCs. The semi-circular

plate model used in the indentation work by Wang et al [82] will then be applied to

approximate the interfacial fracture toughness values.

This study intends to clarify the effects of bond coat and substrate on the degradation

of TBCs, by examining the evolution of TBC interfacial adhesion, TGO growth and

rumpling, and the ’ formation in bond coats against isothermal oxidation. A

systematic analysis was carried out on five different industrial grade TBCs with the

yttria-stabilized zirconia (YSZ) top coat produced by the electron beam physical

vapour deposition (EB-PVD) method on top of 3 different bond coats and substrates,

which then allowed us to establish a relationship between TBC adhesion, TGO growth

and rumpling, bond coat phase distribution. The results have clearly indicated the

83

dynamics (i.e. dependence) of oxide-bond coat interfacial adhesion with respect to the

phase distribution of the bond coats and TGO growth rate. Moreover, it has been

clearly demonstrated that accurate prediction of TBC spallation lifetime could only be

achieved by taking into account the influence of substrate and bond coat chemistry on

the bond coat-TGO interfacial adhesion.

4.2 Experimental details

Three different industrial standard single-crystal nickel based superalloys were used

for the following work: CMSX-4, SRR99, and TMS-82+. The nominal chemical

compositions for these three alloys are given in Table 4.1. These fully heat treated

(solution and primary aged) alloys were processed by conventional investment casting

methods into cylindrical rods (10mm diameter), having the long axis aligned closely

to the {0 0 1} direction. Each rod is further sliced into disk shape (10mm diameter

and 4mm thickness) and spot welded onto Nimonic sticks. On the other hand, three

different bond coats, namely HT Pt-Al (where HT implies high temperature and low

activity aluminium environment), LT Pt-Al; where LT implies low temperature and

high activity aluminium environment), and Pt-diffusion, were subsequently deposited

onto the surfaces of the disk shape alloys. The CMSX-4 and SRR99 specimens were

only paired with LT Pt-Al bond coat for the purpose of studying the influence from

superalloys alone. Table 4.2 below provides a summary of the five different TBC

systems used for this work.

84

Table 4.1

Nominal composition of the three superalloys used in this work (wt%)

Substrat

e

Co Cr Mo W Al Ti Ta Hf Re Ru C Ni

SRR99 5.0 8.0 - 9.5 5.5 2.2 2.8 - - - - Bal.

TMS-82

+

7.8 4.9 1.9 8.7 5.3 0.5 6.0 0.1 2.4 - - Bal.

CMSX-

4

9.6 6.5 0.6 6.4 5.6 1.0 6.5 0.1 3 - - Bal.

Table 4.2

Visual representation of the five TBC systems used in this work

SUPERALLOY SUBSTRATE

BOND COAT CMSX-4 TMS-82+ SRR99

HT Pt-Al ✔

LT-Pt-Al ✔ ✔ ✔

Pt diffusion ✔

These commercially available TBCs, fabricated by Rolls-Royce Plc and High

Temperature Materials Unit of NIMS Japan, were specially selected to cover different

generations of superalloys (SRR99: 1st generation; CMSX-4: 2

nd generation;

TMS-82+: 2nd

generation). The first generation superalloys typically contain more Ti,

but less Re and Ru, while the latest superalloys tend to be Ti-free with higher Re and

Ru. These alloys were then paired with three common bond coat systems: Pt-diffusion,

LT Pt-Al, and HT Pt-Al. The three bond coat systems tend to form very different

concentration profiles in the bond coat region (i.e. Al and Pt contents), as a result of

the different high temperature diffusion processes used during coating fabrication.

The diversity in substrates and bond coats (i.e. difference in compositions) helps to

identify different failure/degradation characteristics. Figure 4.1 below indicates the

85

spallation lifetime of various TBC systems [26], including the substrates and bond

coats used in this work. It can be seen that the CMSX-4 and TMS-82+ have relatively

better spallation lifetime performance than the SRR99 alloy when paired with any of

the three bond coats. On the other hand, the Pt-diffusion (LCBC) bond coat tends to

have the best spallation lifetime performance when paired with any alloy, followed by

the LT Pt-Al (RT22LT), while the HT Pt-Al (CN91PA) comes last in terms of

spallation lifetime.

Fig. 4.1. A comparison of the relative spallation lifetime of commercial TBCs

subjected to thermal cycling, where the x-axis indicates the types of superalloy while

legends indicate bond coat (LCBC = Pt-diffusion, RT22LT = LT Pt-Al, CN91PA = HT

Pt-Al) [26].

The HT Pt-Al and LT Pt-Al bond coat specimens were processed by first

electrodepositing a thin layer of platinum of 5 and 7 m, respectively, followed by a

vacuum heat treatment at 1100°C for 1hr. Next, 5 hrs vapour phase, and 20 hrs pack,

86

aluminization processes were applied at 1080°C and 870°C, respectively. The

Pt-diffusion bond coat was processed similarly by first electroplating a 10 m layer of

platinum, and subsequently treated to a vacuum treatment process at 1150°C for 1hr

without the aluminization process. A further heat treatment in argon atmosphere was

applied on all specimens for 1hr at 1100°C, prior to the deposition with a ZrO2/7wt.%

Y2O3 (YSZ) top coat of 175m by EB-PVD. Lastly, a vacuum heat treatment (1100°C,

1hr) and ageing (870°C, 16hrs) were applied to all specimens prior to testing for

quality assurance purposes.

Five specimens for each TBC system in a total of 25 sliced pieces were subjected to

isothermal exposure at a furnace temperature of 1135°C for 0, 30, 50, 100, and 200hrs.

According to previous studies [37,39], thermal cycling of TBCs leads to significant

interfacial rumpling upon early oxidation times, making it difficult to achieve a

smooth planar interface along the indentation orientation at the interface. Therefore,

selection of isothermal exposure instead of cycling exposure was to be able to study

the influence of the inherent bond coat and substrate chemistry on interfacial adhesion

of the TBCs after thermal exposures.

The exposed specimens were then mounted in a Struer Citovac vacuum impregnation

chamber using a low-shrinkage epoxy resin in order to preserve the specimen edges,

and polished to mirror finish for microstructural characterization. A high resolution

KEYENCE VHX-1000 optical microscope was used to analyze the microstructural

changes on the entire cross-section of all the specimens. Micrographs of the entire

TGO/bond coat interfacial area were recorded for every specimen.

87

Cross-sectional interfacial adhesion test was carried on the thermally exposed

specimens by Vickers micro-indentation based on similar techniques proposed by

previous researchers [51,79,80], using an AMT series MMT-7 Micro-Vickers hardness

machine. The adhesion results were then used to elucidate the mechanisms underlying

the TGO/bond coat interfacial degradation. The symmetrical Vickers indenter was

placed in a way such that one of the diagonal of the indent is as close to the

TGO/bond coat interface as possible (see Fig. 4.2). For an indent to be considered

accurately placed, the entire diagonal needs to lie within a distance of approximately

3m from the TGO/bond coat interface. Indents that did not meet this standard were

discarded from the averaging of crack lengths. Although no systematic study was

carried out regarding the sensitivity of the test results to accuracy of hitting the

interface, it can be seem later in section 4.3.2 that the standard deviations of the crack

length results were generally low enough for a reasonable comparison of the

specimens. Cracks were deliberately introduced at the interface, and the crack length

was calculated by measuring the crack propagation from both corners of the indent.

The length is defined by the total distance of the cracks initiated, starting from either

of the two corners of the indent to the points where the cracks terminate by a distance

of at least 10m (See Fig 4.3 and 4.4). For every specimen, a minimum of 6

indentations was conducted along the interface in order to obtain an overall average

crack length to represent the adhesion of that particular specimen.

88

Fig. 4.2. SEM image showing the placement of the Vickers indenter at the TGO/bond

coat interface.

Fig. 4.3. SEM image showing the crack propagation to the left of the indent.

89

Fig. 4.4. SEM image showing the end of crack propagation to the left of the indent.

Many test trials using different indentation loads were carried out on as coated TBC

samples (i.e. without oxidation). It was found that a load of at least 9800mN (peak

holding time of 10s) was necessary to generate cracks that were visible under SEM.

For the oxidized samples, particularly those treated to longer times, such a load

tended to generate very long cracks, making it difficult to make at least 6 indents per

specimen over the sample cross-section. Hence, all of the isothermally exposed

specimens (i.e. 30-200hrs) were indented with a lower load of 4900mN and a peak

holding time of 10s. Figure 4.5 shows a comparison of the crack lengths of the same

TBC system generated with the two different loads. The crack lengths of the

specimens differed by approximately two folds when the loading doubled. It should

be noted; however, that not all specimens or instances of isothermal exposure follow

this doubling trend.

90

Fig. 4.5. A comparison of typical crack lengths between those generated by 9.8 and

4.5N indentation loads, showing differences of roughly two-folds in length.

The entire coating cross-section of every indented specimens was observed and

photographed using a FEG-SEM at an accelerating voltage of 20kV and a working

distance of 10mm. In addition, compositional mapping for all the major alloying

elements was carried out on the 30hrs specimens near the TGO/bond coat interfacial

region at 10,000X magnification. An additional chemical mapping analysis was done

using EPMA (Electron Probe Microscope Analysis) for the SRR99 LT Pt-Al specimen

under an accelerating voltage of 15kV and a beam current of 20nA.

Residual stress analysis was conducted on all five TBC systems with a set of newly

sliced specimens, which were isothermally exposed at 1135°C for 30, 50, 80 and

100hrs. The residual stress was measured using the photo-luminescence

piezo-spectroscopy method [29,30,104,105], by acquiring Cr3+

luminescence spectra

from -Al2O3. This experiment was carried out using a Renishaw Raman spectrometer

fitted with an optical microscope (Renishaw 1000 Ramanscope system UK). A green

91

argon laser with a wavelength of 514.5 nm was used as the laser source. The spectra

were taken from at least 5 different locations away from the sample edge, each of

which had an acquisition time of 30s. During the course of experiment, the Raman

machine was calibrated once prior to testing a specimen by taking a spectrum from a

strain-free single crystal sapphire sample. The residual stress in the TGO was

estimated from the R2 peak shift given off from the TGO relative to the unstrained

sapphire, assuming a planar equibiaxial stress: (GPa) = Δv (cm-1

) /5.07 [25].

Measurement of the interfacial rumpling was carried out by tracking the TGO/bond

coat interface with three cross-sectional SEM images over a total distance of 573 m.

Each position along the interface was represented by a pair of x (in-plane with the

specimen’s cross-sectional surface) and y (out-of-plane direction parallel to EB-PVD

coating growth direction) coordinates, and that the displacement between two points

was approximated by a straight line.

4.3 Results

4.3.1 Microscopy observation

Fig. 4.6 shows some representative micrographs of three different bond coats (LT

Pt-Al, HT Pt-Al and Pt diffusion) paired with TMS-82+ superalloys. According to

these micrographs, TGO thickness, rumpling and precipitates can be observed and

compared. For example, the TMS-82+ LT Pt-Al specimen seemed to have slightly

thicker TGO growth and noticeably more rumpling than that of the TMS-82+ HT

Pt-Al at 30hrs. At 50hrs, both specimens began to form large pieces of precipitation

92

near the TGO/bond coat interfacial region. Such precipitates have identified as ’

phase formed after thermal treatments. Fig. 4.7 shows representative images of LT

Pt-Al bond coat with CMSX-4, SRR99, TMS-82+ superalloys as substrate, where the

SRR99 specimen has a much thicker TGO growth than that of the TMS-82+ and

CMSX-4. At 30hrs, large pieces of ’ phase could already be seen in the SRR99 LT

Pt-Al specimen, yet both the CMSX-4 and TMS-82LT Pt-Al specimens had only trace

amount of those near the grit line further away from the interfacial region.

93

Fig. 4.6. Representative interfacial microstructure between the three different bond

coat systems (LT Pt-Al, HT Pt-Al, Pt-diffusion) paired with TMS-82+ superalloy. The

HT Pt-Al bond coat had the least rumpling, yet failed to survive up to 200hrs.

94

Fig. 4.7. Representative interfacial microstructure between the three different

superalloys (CMSX-4, TMS-82+, SRR99) paired with LT Pt-Al bond coat (SRR99

failed after 30hrs, and no photos are available for 100 and 200 hrs time trials). The

SRR99 specimen had the most ’phase formation and rumpling, and could not survive

95

up to 50hrs.

4.3.2 Lifetime and interfacial adhesion of TBC systems

A qualitative analysis of the TGO/bond coat microstructure is given in this section for

the three superalloys and bond coats. First, the thermally treated specimens were

examined by optical microscopy for whether or not spallation has occurred (defined

as approximately 20% detachment of the top coat by cross-sectional length). It was

found that the SRR99 LT Pt-Al specimen was amongst the first to fail between 30 and

50 hrs, as its coating remained mostly intact at 30hrs, but major spallation was soon

observed at 50hrs. The TMS-82+ HT Pt-Al also had reached its spallation life

between 100 and 200 hrs, as indicated by over 20% of coating detachment at 200hrs.

All remaining specimens survived well into 200hrs exposure without noticeable

coating detachment at all.

A quantitative analysis is done using the cross-sectional indentation technique to rank

the interfacial adhesion for the three different substrates and bond coats. Fig. 4.8

shows the crack lengths against thermal exposure time for the five different TBC

systems. CMSX-4 with LT Pt-Al bond coat with lifetime longer than 200hrs exhibited

the shortest crack lengths of all specimens, having lengths no greater than 50m

throughout the entire thermal exposure time. SRR99 LT Pt-Al with lifetime between

30-50hrs, on the other hand, had the longest cracks after only 30hrs thermal exposure

with an average length of 128m. The three TMS-82+ specimens had a somewhat

intermediate crack resistance, with the HT Pt-Al bond coat having the longest average

96

crack length of 128 m after 100hrs thermal exposure.

Fig. 4.8. The variation of crack lengths vs. isothermal oxidation (1135°C) hours for the

five different TBC systems (i.e. the prefix TMS-82+ is substrate, while Pt-diffusion is

bond coat) using a load of 4900mN.

It was noticed that the CMSX-4 LT Pt-Al and TMS-82+ HT Pt-Al specimens did not

follow a typical increasing crack length trend with the number of exposure hours. The

CMSX-4 specimen, for instance, had a declining length trend going from 44 m at

30hrs to 17 m at 100hrs, before rising to 38 m at 200hrs. The HT Pt-Al specimen

also had similar behavior, for which its crack length went from 111 m at 30hrs to 36

m at 50hrs, before rising to 128 m at 100hrs (failure after 100hrs). This

phenomenon suggests the need for further investigation on these two specimens, by

97

studying their interfacial evolution, and by comparing them with other specimens.

4.3.3 TGO thicknesses, residual stresses and rumpling at the TGO/bond coat

interface

Fig. 4.9 shows the oxidation kinetics (average TGO thickness vs. time1/2

) curve. First,

it can be seen that the TGO growth rate started out quite similar for all (ranging

approximately from 0.0806 to 0.0830 m/hr), except the SRR99 LT Pt-Al specimen

which had a higher growth rate of over 0.100 m/hr. This specimen was not able to

survive up to 50hours. As the exposure time increased, the variation in thicknesses

between different specimens became more apparent, particularly for the comparison

between CMSX-4 LT Pt-Al and other specimens. The CMSX-4 LT Pt-Al specimen

had maintained a very thin TGO layer, and that its growth rate became relatively slow

between 30 and 100hrs at only 0.00572m/hr, whereas the 2nd

slowest specimen has

a growth rate of 0.0137m/hr. As shown earlier, CMSX-4 LT Pt-Al had the best

interfacial adhesion, and a lifetime longer than 200hrs.

98

Fig. 4.9. The TGO growth kinetics vs. isothermal oxidation (1135ºC) hours for the five

different TBC systems.

Fig. 4.10 shows the residual stress in the TGO of every specimen from 30 hr to 100hrs,

obtained using luminescence measurements. SRR99 LT Pt-Al specimen shows the

lowest stresses in the TGO which spalled after 30hrs. Among the other 4 samples, the

stress level in TGO of both TMS-82+ HT Pt-Al and CMSX-4 LT Pt-Al shows

decreasing trend with increase in the thermal treatment time, while the other 2

samples show slight increase of the stresses with increase in the thermal exposure

time.

99

Fig. 4.10. The TGO compressive stress vs. isothermal oxidation (1135°C) hours for

the five different TBC systems.

Fig. 4.11 shows the evolution of TGO length against thermal exposure time for up to

100hrs. It can be seen that the initial degree of rumpling was very similar for different

samples, but a noticeable difference was soon observed after 30 hrs of thermal

exposure. TMS-82+ LT Pt-Al had the most rumpling through the course of

experiment up to 100hrs, while the TMS-82+ HT Pt-Al specimen maintained the least

rumpling throughout the entire exposure time. On the other hand, the TMS-82+

Pt-diffusion specimen did not show much increase in rumpling throughout the entire

exposure time.

100

Fig. 4.11. The evolution of interfacial rumpling vs. isothermal oxidation (1135°C)

hours for the five different TBC systems.

4.3.4 Phase transformation and Ni/Ti segregation

The precipitation mentioned previously is identified as to ’ phase transformation,

which according to the Ni-Al-Pt ternary phase diagram in the work of Gleeson et al

[36], occurs when Al is undergoing depletion as a result of oxidation. Quantification

of this phase was carried out by measuring the total amount by area (m2) of the

precipitates that are attached to the TGO/bond interface over a distance of

approximately m. This quantification was done using the aid of an image

analysis program. Fig. 4.12 shows the content of the ’ phase as function of thermal

treatment time. The SRR99 LT Pt-Al specimen had the fastest ’ formation in the

101

early 30hrs with the shortest lifetime. The 2nd

most ’formation can be seen on the

TMS-82+ HT Pt-Al specimen, followed by TMS-82+ LT Pt-Al and CMSX-4 LT

Pt-Al (see Fig 4.12).

Fig. 4.12. The amount of ’ formation attached to the TGO/bond coat interface vs.

isothermal oxidation (1135°C ) hours.

EDX analysis was done on all specimens near the TGO/bond coat interfacial area. Fig.

4.13 shows the Ti maps while other elemental maps either had no obvious variation

between them, or were slightly influenced by the dilution effect from the

aluminization process. A noticeable level of Ti concentration was present near the

bond coat of the SRR99 LT Pt-Al system, including a patch of highly concentrated

area near the interface. This was however, much less obvious in the case of the

102

TMS-82+ and CMSX-4 LT Pt-Al systems. The observed Ti segregation phenomenon

on SRR99 LT Pt-Al was further examined using EPMA, and the results are shown in

Figure 4.14. A clear segregation between Al/Cr and Ni/Ti could be seen near the

TGO/bond coat interfacial area.

Fig. 4.13. EDX mapping showing Ti distribution near the TGO/bond coat interface of

the coating after 30hrs of isothermal exposure. An obvious Ti enrichment can be seen

in the bond coat of the SRR99 specimen.

103

Fig. 4.14. EPMA mapping showing the elemental mapping of Ni, Ti, Al, and O for

SRR99 LT Pt-Al after 30hrs of isothermal exposure (the color bars are values in wt

%). The regions of Al and Cr depletion correspond exactly to the enrichment in Ni

and Ti contents.

As shown in the previous sections, the phenomena of early to ’phase

transformation and Ti segregation were both found near the TGO/bond coat region of

the SRR99 LT Pt-Al specimen. Hence, there is an interest to investigate whether or

not there is any link between the two phenomena. A thermodynamic software, named

Thermo-Calc was used to identify Ti distribution in the and ’ phases of the Pt-Al

bond coat system. A set of conditions including the isothermal temperature used

(1409K), the oxidizing atmospheric pressure (1 atm), the Al concentration of

as-coated Pt-Al bond coat at the interface (19 wt%), and the overall composition of

the major alloying elements (Ni balanced) were specified as the input for the

104

calculation. At any given Al concentration, we assume that the system is in a stable

equilibrium state after certain amount of time. Note that the 19 wt% Al was chosen

based on the EDX analysis result of the SRR99 LT Pt-Al sample studied by Wu et al

[26]. With increase of the Al content, all other remaining alloying elements were

proportionally reduced based on their original wt%, so as to reflect the dilution effect

from the aluminization process (see Table 4.3).

Table 4.3

The nominal composition in wt% of the bond coat of the SRR99 specimen as

specified in the calculation of Thermo-Calc relative to its original bulk composition

Substrate Co Cr Mo W Al Ti Ta Hf Re Ru C Ni

SRR99 5.0 8.0 - 9.5 5.5 2.2 2.8 - - - - Bal.

SRR99 (bond

coat)

4.3 6.9 - 8.2 19.

0

1.9 2.4 - - - - Bal.

Fig. 4.15a gives the mole-fraction of different phases in SRR99 LT Pt-Al obtained

based on the calculation in Thermo-Calc, starting from 19 wt% Al and ending at 10

wt%. This range is set in order to simulate Al depletion due to isothermal oxidation of

the experimental specimen. The Ti distribution results in the and ’ phases, were

calculated with 12% Al concentration, ensuring that both phases are present on the

three alloys compared (note that the relative Ti distribution trend does not vary with

Al%). Fig. 4.15b shows the mole fraction of Ti in each of the two phases for the three

different alloy specimens. The calculation results suggest that the depletion of Al

leads to formation of the ’ phase while the Ti enriched in the ’ phase. It should be

noted that the database that came with this version of Thermo-calc here does not

contain the Pt element. Hence, the thermodynamic results were calculated under the

assumption that the Pt content is substituted entirely by Ni element. Since Pt and Ni

105

belongs to same platinum group metals with similar chemical and physical properties,

these elements tend to mix together in solid solution (i.e.,(Pt,Ni)Al of the -NiAl

bond coat phase). Hence, it is reasonable to assume that substitution of Pt by Ni in the

calculation does not influence the phase formation results obtained above.

Fig. 4.15. a) The mole-fraction of different phases existing as a function of Al content

in the SRR99 LT Pt-Al sample obtained from calculation using Thermo-calc, and b)

The comparison of Ti concentration in and ’ phases of the three alloys.

4.4 Discussion

4.4.1 Approximation of interfacial fracture toughness Gc

As mentioned in Chapter 2, it is desirable to obtain a quantitative comparison of the

adhesion between any two coating systems based on experimental results, in

particular, microstructure features generated by the cross-section indentation. Out of

all the techniques introduced in Chapter 2, the semi-circular clamped model using

106

cross-sectional indentation as the mean to generate fracture appears to be applicable to

produce reliable fracture toughness values of TBCs without the hassle of complicated

sample preparation and loading setup [77]. Hence, it is utilized here to quantify the

interfacial fracture toughness of the five TBC systems, which can be represented by

the energy release rate Gc with respect to crack radius, a, as shown in Eq.4-1 below:

(Eq. 4-1)

It should be noted that the indentation in this work generated mostly out-of-plane

stresses (i.e. opening tensile force at the interface) without buckling in the

delaminated coating. Therefore, it is reasonable to neglect the influence of in-plane

stresses (i.e. stored residual stress in the coatings), which can contribute to additional

energy release. It can be noted that the crack length generated by indentation is much

less than the critical spallation size by buckling, eg. a critical buckling crack radius of

886m for the 130m TBC coating in Wang’s work [86]. The values of Ec are

approximated from the data plot in Fig. 6 of the work on YSZ stiffness using

miniature 3 point-bending by Wang et al [106]. Since the bi-layer coatings in that

work consisted of 8 wt% YSZ and a TGO scale formed on a Pt-diffusion bond coat,

the effective modulus obtained should bear close resemblance to that of the coatings

used in this work as similar conditions were used to thermal treat the TBCs (eg.

EB-PVD was applied at 1000°C, while furnace temperature of 1150°C was used for

thermal cycling). The coating thickness hc is taken as the sum of YSZ thickness

(175m) and TGO thickness (given in Figure 4.9). The Poisson’s ratio is taken as

0.3 is this case. The values of Ec and hc are summarized below for 0, 30, 50, and

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100hrs isothermal exposures in Table 4.4 and 4.5, respectively. Figures 4.16 shows the

crack radius of the five TBCs isothermally exposed to 0, 30, 50, and 100hrs. Note that

the crack radius is calculated as the sum of the crack lengths (given in Figure 4.8) and

half-diagonal of the indent marks. The central out-of-plane displacement, u, cannot be

measured accurately, since the displacement consists of both elastic and plastic

deformation. Hence, it was first approximated as the half-diagonal of the indentation

mark. The calculated Gc values, however, were many orders of magnitude higher than

the values given in previous literatures [54,82] (Table 4.6), thus suggesting that

half-diagonal of the indent mark is an overestimation for u.

Table 4.4

The effective in-plane Young’s modulus of the bi-layer coating (YSZ-TGO) after

different isothermal exposures (These values were estimated from Fig. 6 in Ref.[106])

Isothermal hrs 0 30 50 100

Ec (GPa) 7.5 13 18.5 13.5

Table 4.5

The coating thicknesses (YSZ+TGO thickness) values used for the calculation of Gc

in this work

hc (m)

Isothermal hrs 0 30 50 100

CMSX-4

LT Pt-Al 175 177.42 177.73 177.82

TMS-82+

LT Pt-Al 175 177.49 177.81 179.34

108

TMS-82+

HT Pt-Al 175 177.44 177.74 178.92

TMS-82+

Pt-diffusion 175 177.45 177.54 178.40

SRR99

LT Pt-Al 175 178.15 -- --

Table 4.6.

A comparison between the Gc reported in Ref. [82] and the TMS-82+ Pt-diffusion

here by approximating u as half-diagonal of the indent mark.

Gc (J/m-2

)

Isothermal hrs 0 30 50 100

TMS-82+

Pt-diffusion 2.5x10

8 2.7x10

6 1.2x10

6 1.2x10

6

Thermal cycles

(1-hr cycle) 2 30 50 100

CMSX-4

Pt-diffusion 2300 80 42 42

109

Fig. 4.16. The crack radius (i.e. sum of the length of half indent and crack lengths)

values used in the calculation of Gc in this work.

In order to get a more reasonable value of Gc, u value should be much smaller than

the half-diagonal of the indent. From SEM observation, the diamond shape indent

caused mostly inward displacement (i.e. plastic deformation) of the YSZ with

minimal central out-of plane (i.e. parallel to YSZ columns) deformation. By applying

an opening tensile force at the TGO/bond coat interface using indentation, there is

very little upward movement in the YSZ top coat due to its porous/soft structure.

However, the bond coat likely displaced downward, for which the displacement u can

be roughly approximated as the crack openings at the TGO/bond coat interface near

the two corners of the indent mark (Fig. 4.17). 6 measurements (i.e. 3 from each of

the corner parallel to interface) were taken from a single indent, and used for the

110

averaging of upward displacements. Figure 4.18 shows the u values of the five TBCs

isothermally exposed to 0, 30, 50, and 100hrs. Finally, the interfacial fracture

toughness is measured by using all the above attributes and summarized in Figure

4.19. There seems to be a wide scattering of Gc values between the five TBCs, which

follow a similar trend to that of the crack length results shown earlier. The Gc values

given here for the as-coated TBCs (i.e. 0hr) seem to be much lower than those

reported in the work by Wang et al [82]. In that work, the Gc for their CMSX-4

Pt-diffusion specimens exposed to only 2 thermal cycles had a much larger value

(2300Jm-2

). It is possible that the u values (taken as the crack openings due to

downward displacement of bond coat) for the as-coated TBCs were significantly

underestimated for the calculation of Gc in the case here. Despite using a load two

times greater (9.8N), the u values (i.e. crack openings) are roughly the same if not

lower than those of later isothermal hours (i.e. 30-200hrs), which were indented with

4.5N. As shown in Figure 4.20, there seems to be significantly more damage in the

as-coated YSZ in the form of inter-columnar fracture and shear displacement than the

same coating after 30 hrs of isothermal exposure. This can be explained by the

sintering of the YSZ top coat after prolonged thermal exposure, which prevented the

stress generated by the indentation from dissipating in the form of YSZ damage in

isothermally exposed TBCs.

111

Fig. 4.17. Micrograph showing the location of the u measurements (6 per indent).

Fig. 4.18. The upward displacements (i.e. taken as the crack openings at the

TGO/bond coat interface near the corner of indents) used in the calculation of Gc in

this work.

112

Fig. 4.19. The approximated values of Gc for the five TBCs based on the semi-cirular

plate model.

113

Fig. 4.20. Micrographs showing the extent of YSZ damage between as-coated TBC

and TBC subjected to 30hrs of isothermal exposure.

Other uncertainties may also contribute to the discrepancy between the Gc values

obtained here and those given by other literatures. For instance, the Ec values taken

from Wang’s work [106] were measured on specimens that underwent thermal cycling

for the same amount of time instead of isothermal exposure as in this work. There

would certainly be a difference in the Ec values between isothermal exposed and

thermal cycled TBCs of the same duration, due to different amount of top coat

sintering and residual stresses present in the top coat and TGO. In addition,

differences in coating deposition methods, coating thicknesses, bond coat structural

change, and phase transformation may also be sensible reasons for the discrepancy in

the Gc values, as these differences can change any of the inputs in Eq. 4-1.

Nevertheless, the interfacial fracture toughness values calculated here for the

TMS-82+ Pt-diffusion are fairly similar to that of the CMSX-4 Pt-diffusion specimen

114

reported in Wang’s work [82]. The Gc values obtained for the 30, 50, and 100hrs

isothermally exposed TMS-82+ Pt-diffusion specimens are 62, 31, and 26 J/m-2

,

respectively, which are fairly close to the values obtained by Wang as shown in Table

4.6.

4.4.2 Linear fitting of crack rate with degradation factors

In order to compare interfacial crack resistance under indentation for different

samples, linear fitting has been done on crack length versus thermal exposure time

(hrs) and then the slope of each fitting has been defined as the crack rate. A higher

slope means a higher crack tendency. Therefore, it should be noted that the crack rate

(m/hr) is not a measure of the crack propagation speed under indentation, instead, is

used for the comparison of crack tendency among different samples under indentation.

In addition, TGO thickness growth rate, TGO length growth rate (rumpling) and

’phase formation rate are obtained based on plots in Fig. 4.9, 4.11 and 4.12 where

’phase formation rate was obtained based on initial treatment time only. As the TGO

stresses do not vary significantly during thermal treatments (Fig. 4.10), therefore

average stress value for each sample during the thermal treatment period is taken for

comparison of TGO stresses in different samples. It should be noted again here that

the purpose to define these parameters is not to obtain absolute value of each

parameter, instead, it is to compare properties of different samples, i.e. to identify

effect of different factors on crack resistance and lifetime of TBCs. Table 4.7 gives

ranking of crack rate, TGO thickness growth rate, TGO length growth rate (rumpling),

TGO stresses and ’phase formation rate for the 5 different samples. In addition,

quantitative values for these parameters are also given in histograms for the 5

115

different samples (Fig. 4.21-4.25).

Table 4.7

The comparison between the rankings of performance and contribution factors studied,

showing an exact match between performance and the ’ formation factor

Performance Contributing Factors

Best(Factors:

low)

Worst

(Factors:

high)

Crack rate TGO growth

rate

TGO Stress TGO

rumpling rate

’ formation

rate

CMSX-4 LT

Pt-Al

CMSX-4 LT

Pt-Al

SRR99 LT

Pt-Al

TMS-82+

HT Pt-Al

CMSX-4 LT

Pt-Al

TMS-82+ LT

Pt-Al

TMS-82+ Pt

diffusion

TMS-82+ LT

Pt-Al

TMS-82+ Pt

diffusion

TMS-82+ LT

Pt-Al

TMS-82+

Pt-diffusion

TMS-82+

HT Pt-Al

CMSX-4 LT

Pt-Al

CMSX-4 LT

Pt-Al

N/A

TMS-82+

HT Pt-Al

TMS-82+ LT

Pt-Al

TMS-82+

HT Pt-Al

TMS-82+ LT

Pt-Al

TMS-82+ HT

Pt-Al

SRR99 LT

Pt-Al

SRR99 LT

Pt-Al

TMS-82+ Pt

diffusion

SRR99 LT

Pt-Al

SRR99 LT

Pt-Al

116

Fig. 4.21. A quantitative comparison of crack rates amongst the five TBC specimens.

Fig. 4.22. A quantitative comparison of TGO growth rates amongst the five TBC

specimens.

117

Fig. 4.23. A quantitative comparison of average TGO stresses amongst the five TBC

specimens.

Fig. 4.24. A quantitative comparison of TGO rumpling rates amongst the five TBC

118

specimens.

Fig. 4.25. A quantitative comparison of ’phase formation rates amongst the five TBC

specimens.

Fig. 4.26 gives the crack rate as function of ’phase formation rate, TGO thickness

growth rate, TGO length growth rate (rumpling rate), and TGO average stress. Both

the ’phase formation rate and TGO thickness growth rate appear to correlate with the

crack rate (i.e failure of TBCs) in a linear relationship. In contrast, the TGO stress

shows no obvious correlation with the crack rate. Finally, the TGO rumpling

measured based calculating the growth rate of TGO length (i.e. rumpling rate) also

has no apparent correlation to TBC failure, which appears to contradict previous

studies [107,108].

119

Fig. 4.26. Plots showing the linear dependency of crack rate on a) ’phase formation

rate, b) TGO thickness growth rate, c) TGO length growth rate (rumpling rate) and, d)

TGO average stress. The crack rate only increases linearly with the rate of ’phase

formation.

4.4.3 Possible degradation mechanisms by ’ formation rate

Although an obvious correlation can be seen between the ’ formation rate and the

crack rate, there is no clear explanation regarding how the early ’ formation may

120

have contributed to the loss of interfacial adhesion at this point. It is speculated that

the presence of ’ phase (rich in Ni and Ti) near the interface could either hinder the

formation of alumina, or may alternatively form titanium oxide. These observations

suggest that the two weakest samples: SRR99 LT Pt-Al and TMS-82+ HT Pt-Al

specimens had undergone a chemical diffusion that was quite different from other

specimens, and that the formation of Ni and Ti enriched ’ phase either directly or

indirectly reduced the TGO adhesion. Further chemical diffusion analysis is necessary

to fully understand the effect of Ni and Ti-containing ’ phase on the interfacial

adhesion.

It has been known that fast growing Co and Ni-rich oxides could form as TGO aside

from the alumina growth, and are more pronounced in specimens with lower

aluminum concentration profile in their bond coats. This failure mechanism may have

been closely linked to the early ’ formation, which is reflected by the adjusted

r-square values in the linear fit of Fig. 4.26.

4.4.4 Stress relaxation by rumpling

A surprising phenomenon was observed from the crack length measurements for

CMSX-4 LT Pt-Al and TMS-82+ HT Pt-Al, of which fluctuation in crack length was

found. The residual stress shown earlier indicated that both the CMSX-4 LT Pt-Al and

TMS-82+ HT Pt-Al specimens underwent a stress decline from 30 to 100 hrs and 30

to 80 hrs respectively (Fig. 4.10). Interestingly, declining trends in crack length were

also found at the same time intervals on those two specimens. This phenomenon was

not observed in the other three specimens, for which their corresponding crack length

121

all increased with thermal exposure. Further, quantification of the bond coat/TGO

interfacial rumpling showed that CMSX-4 LT Pt-Al and TMS-82+ HT Pt-Al are the

only specimens undergoing further rumpling between 30 and 50 hrs. By taking into

account this increasing rumpling behaviour with the corresponding stress relaxation

shown earlier, it is believed that the observed fluctuation in crack length may have

been due to the relaxation of TGO residual stress by rumpling [28,30,31].

Fig. 4.27 shows that average TGO stress reduces with increase in rumpling rate,

which suggests that TGO rumpling can relax TGO residual stress. It has been shown

by geometric factors that a greater out-of plane tensile stress tends to accumulate at

the protuberance of the TGO/bond-coat interface, and at the YSZ side of the interface

along the flanks of the protuberance for specimens with rougher interface [109,34].

This claim is supported by the fact that specimens with increased interfacial

roughness tend to fail not only at TGO/bond coat, but also at TGO/YSZ interfaces

[83]. While such an introduction of stresses along the interfaces is undoubtedly

detrimental, our experiment results have shown several cases of a specimen with

higher degree of rumpling, yet its interfacial adhesion and lifetime were much better

relative to those with less rumpling.

122

Fig. 4.27. The dependency of stress on the rumpling increase rate (highly linear

dependent).

Traditionally, both residual stress and rumpling at the TBC interface were considered

as important factors attributing to the TBC failure. However, in this study, a direct

correlation could not be established between these two factors and the TBC spallation

life (Fig. 4.26). The CMSX-4 LT Pt-Al specimen has been shown to possess the best

crack resistance amongst all the specimens, despite having a moderate TGO stress and

a moderate rumpling rate (Table 4.7). The reason behind this excellent performance

may be largely due to a slow ’ formation. Up to 100hrs, the CMSX-4 specimen had a

marginally less ’ formation than both the TMS-82+ LT Pt-Al and HT Pt-Al

specimens. Beside the kinetics of ’ formation, the TGO growth rate also seemed to

bear some influence on the crack resistance. Throughout the entire thermal exposure,

CMSX-4 specimens demonstrated the slowest TGO growth kinetics relative to

specimens of other substrates, although the difference was much less obvious than in

123

the case of ’ formation.

4.4.5 The dynamic feature of interfacial fracture toughness

One key implication of this work is that the TBC spallation lifetime is dependent on

both the variation of driving force G for spallation and interfacial fracture toughness

Gc. This concept had been initially proposed by Wu et al [35], arguing that the focus

of TBC modeling had been largely placed on the origin of the driving force for

spallation, without thoroughly considering the dynamic nature (i.e. variation) of the

interfacial fracture toughness with respect to the substrate chemistry. In that original

work, this dynamic nature was indirectly demonstrated based on preliminary

observation of the TBC spallation lifetime. The current work, in contrast, has clearly

demonstrated the sensitivity of the interfacial fracture toughness to substrates and

bond coat systems. This can be illustrated in Fig 4.28, a semi-quantitative plot of the

variation of the driving force G (steady state) for spallation and interfacial fracture

toughness Gc of TBC specimens between 30 and 100hrs isothermal exposure. The

driving force G, in this case, is calculated by the equation for thin film of thickness h

on a thick substrate and the case of edge delamination [110,111],

(Eq. 4-2)

where E and are the Young’s modulus and Poisson’s ratio of the film, taken as

380GPa and 0.25, respectively. The value is taken as the TGO stress measured in

section 4.3.3. One can see that the TMS-82+ HT Pt-Al having a drastic decline in

interfacial fracture toughness Gc along with its rapid increasing driving force G after

124

80hrs of isothermal exposure, resulted in the early failure of the specimen. On the

other hand, TMS-82+ LT Pt-Al still had vast remaining life as shown by the difference

in energy △G2. Note that the Gc values are roughly approximated by taking the ratios

of crack lengths between the TMS-82+ HT Pt-Al and other specimens at different

times, and multiplying them by the interfacial adhesion value Gc of TMS-82+ HT

Pt-Al, while assuming that the Gc value is 20% higher than its steady state energy

release rate G at 100hrs just prior to failure. Hence, the plot does not necessary

represent the true values obtained using proper testing procedures. Note that the rising

and dropping trend seen in this plot for TMS-82+ HT Pt-Al is due to the fluctuating

crack lengths obtained previously.

Fig. 4.28. A semi-quantitative plot illustrating the variation of driving force G (steady

state) and fracture toughness Gc between 30 and 100 hrs of isothermal exposure. △G

gives an indication of the remaining spallation life in the TBC.

125

4.5 Conclusion

The results of this work have led to the following conclusions:

The cross-sectional indentation technique has been shown as an effective way to

give a semi-quantitative indication of the TBC interfacial adhesion. The crack

propagated mostly along the TGO/bond coat interface, and the results

correspond well with the known spallation lifetimes of commercial TBCs.

Among the 5 TBC samples, the growth rate, rumpling, and residual stress of

TGO as well as the phase distribution of the bond coat have been compared in

relation to the TBC adhesion.

An approximation of the interfacial fracture toughness, Gc, was carried out based

on the semi-circular plate model by Sanchesz et al [112] and the later work by

Wang et al [82]. Some of the calculated results, particularly for the as-coated

TBCs, were many orders of magnitude lower than those reported in previous

work [55,75,82]. Nonetheless, the Gc values for the TMS-82+ Pt-diffusion

specimens in this work (i.e. 30-100hrs isothermally exposed) were determined

to be similar to that of the CMSX-4 Pt-diffusion in Wang’s work [82].

The bond coat phase distribution, i.e. to ’ transformation due to depletion of

Al, appears to be the dominant factor in degrading the interfacial adhesion of

the specimens studied while Ni and Ti concentrate in the ’ phase. In addition,

the TGO growth rate also contributes to the degradation of the TBC adhesion.

It is speculated that the highly concentrated Ni and Ti-containing ’ formation

attached to the interface can either hinder the stable formation of alumina or

alternatively form titanium oxide, which would weaken the TBC interface.

126

The experimental results have shown that interfacial adhesion not only varies

with different materials, but also depends dynamically with the thermal

exposure history of oxidation. However, there is no standard procedure in the

measurement of Gc so far, resulting in large disparity across the values

obtained by different research groups.

127

CHAPTER 5

The degradation mechanisms of TBCs: Driving force for

spallation versus interfacial fracture toughness Gc

5.1 Introduction

For achieving a balance between fuel economy and optimal thrust power in modern

jet engines, there is an ongoing pursue for a higher turbine inlet gas temperature [6].

In order to cope with the ever-demanding operating conditions, state-of-the art

thermal barrier coatings (TBCs) based on yttria-stabilized zirconia (YSZ) have been

widely applied on top of the bulk Ni-based superalloy, to protect the underlying alloy

from approaching its melting temperature. Due to the porous columnar structure and

high ionic diffusivity of YSZ in the ceramic top coat at high temperature, the ingress

of oxygen can readily oxidize the underlying substrate, leading to catastrophic failure

of the turbine structure. To address this issue, an intermediate bond coat layer,

commonly made by either a thermally sprayed MCrAlX (M=Ni or combination of Ni

and Co, and X indicated the minor element addition) or electrodeposition of platinum

aluminide, is applied between the ceramic top coat and the bulk superalloy. The bond

coat contains a relatively higher amount of Al content than the bulk alloy, and hence,

serves as a reservoir of Al to promote the preferential formation of a protective

alumina as the thermally grown oxide (TGO) layer [14,15,16].

Despite these protective measures, the TBCs are still prone to spallation failure

primarily driven by the interfacial rumpling of the bond coat during the course of high

128

temperature oxidation and thermal cycling [90,99,23,24]. The general consensus for

the cause of interfacial rumpling is thought to be the mismatch in the coefficient of

thermal expansion between the multilayered structures of TBCs [23,24]. Despite

considerable progress that has been made on modeling the interfacial evolution of

TBCs in terms of the stress state and the degree of rumpling, there has been a lack of

understanding on how these values could be influenced by the chemical diffusion and

phase transformation of bond coat and substrate during thermal cycling

[23,24,34,103]. It is only in recent years that more and more studies have focused on

comparing and understanding the influence of bond coat and substrate compositions

on the degradation of TBCs [26,35,95,101,102].

As identified in the previous chapter, the rate of to ’ transformation driven by Al

depletion was determined as a dominant factor in degrading the interfacial adhesion of

the TBC samples studied. It was speculated that the ’ formation containing highly

concentrated Ti near the bond coat/TGO interface may weaken the interface by

forming TiO2. However, preliminary analysis by electron probe microscope analysis

(EPMA) indicated that while traces of TiO2 were present in the TGO layer of the

SRR99 LT Pt-Al specimen, the same phenomenon could not be observed on the

TMS-82+ HT Pt-Al specimen, which also had similar Ti enrichment, after the ’

formation near the bond coat/TGO interface (see Figure 5.1). This implies that the

formation of less adhesive oxide cannot be the only mechanism causing the interfacial

degradation of TBCs. Both SRR99 LT Pt-Al and TMS-82+ HT Pt-Al showed very

poor interfacial adhesion according to the results of the previous chapter.

129

Fig. 5.1. EPMA mapping showing the formation of Ti concentrated ’ in the bond coat

region near the oxide-bond interface of the SRR99 LT Pt-Al (containing traces of

TiO2 in the TGO) and TMS-82+ HT Pt-Al after 30 and 100hrs of isothermal exposure,

respectively.

The current study is therefore, devoted to clarifying the difference in the degradation

mechanisms between the SRR99 LT Pt-Al and TMS-82+ HT Pt-Al specimens. A

systematic approach similar to the previous chapter, was carried out by examining

five different industrial grade TBCs based on 3 different bond coats and substrates.

Electron backscatter diffraction (EBSD) pattern maps near the interface were used in

conjunction with high resolution optical micrographs, EPMA chemical mappings, and

nano-indentation hardness values to elucidate the underlying mechanisms behind the

early spallation of TBCs. The results have clearly indicated that the TBC interfaces

may degrade not only by the formation of less adhesive oxides (i.e. TiO2), but also

due to the misfit stress induced in the TGO layer resulting from the mechanical

mismatch between the different phases during cooling. Moreover, the results here

have clearly demonstrated the previously proposed concept that the TBC spallation

130

life is dependent on the dynamic variation of both the driving force for spallation G

and interfacial fracture toughness Gc .

5.2 Experimental details

The same set of TBCs, consisting of three industrial standard single-crystal nickel

based superalloys: CMSX-4, SRR99, and TMS-82+, used in the previous chapter, was

utilized in this current work. Table 5.1 shows the nominal chemical composition for

these alloys. These fully heat treated alloys (solution and primary aged) were

processed using conventional investment casting methods into cylindrical rods (10

mm diameter), having the long axis aligned closely to the {001} direction. These rods

were further sliced into disk shape (10 mm diameter and 4 mm thickness) and spot

welded onto Nimonic sticks. On the other hand, three different bond coats, namely

HT Pt-Al (where HT implies a high temperature and low activity aluminization

process), LT Pt-Al (where LT implies a low temperature and high activity

aluminization process), and Pt-diffusion (without aluminization), were subsequently

deposited onto the flat surfaces of the button-shaped discs. The TMS-82+ alloy was

paired with all three bond coats for the purpose of studying the effect of bond coat,

while CMSX-4 and SRR99 were only paired with LT Pt-Al bond coat to study the

influence from superalloy substrate, when the type of bond coat is fixed. Table 4.2

below provides a summary of the five different TBC systems used in this work.

131

Table 5.1

Nominal compositions of the three superalloys used in this work (wt%)

Substrat

e

Co Cr Mo W Al Ti Ta Hf Re Ru C Ni

SRR99 5.0 8.0 - 9.5 5.5 2.2 2.8 - - - - Bal.

TMS-82

+

7.8 4.9 1.9 8.7 5.3 0.5 6.0 0.1 2.4 - - Bal.

CMSX-

4

9.6 6.5 0.6 6.4 5.6 1.0 6.5 0.1 3.0 - - Bal.

Table 5.2

Visual representation of the five TBC systems used in this work

SUPERALLOY SUBSTRATE

BOND COAT CMSX-4 TMS-82+ SRR99

HT Pt-Al ✔

LT-Pt-Al ✔ ✔ ✔

Pt diffusion ✔

The HT Pt-Al and the LT Pt-Al bond coat specimens were prepared by first

electrodepositing a thin layer of platinum of 5 and 7m, respectively followed by a

vacuum heat treatment at 1100°C for 1 h. Next, 5hrs vapour phase, and 20hrs pack,

aluminization processes were applied at 1080°C and 870°C, respectively. The Pt

diffusion bond coat specimen was first electroplated with a 10m layer of platinum

and then vacuum heat treated at 1150°C for 1hr without the aluminization process. A

further heat treatment in argon atmosphere was applied on all specimens for 1hr at

1100C, prior to the deposition of a ZrO2/7wt% Y2O3 (YSZ) top coat of approximately

175m in thickness by electron beam physical vapor deposition (EB-PVD). Lastly, a

vacuum heat treatment (1100 °C, 1hr) and ageing process (870°C, 16hrs) were applied

132

to all specimens for quality assurance purposes.

A total of 25 specimens (5 for each TBC system) were subjected to isothermal

exposure at a furnace temperature of 1135°C for 0, 30, 50, 100 hrs. These specimens

were then mounted in a StruerCitovac™ vacuum imprenation chamber using a

low-shrinkage epoxy resin in order to preserve the specimen edges, and polished to

mirror finish for microstructural characterization. A high resolution KEYENCE

VHX-1000 optical microscope was used to observe the microstructural features on the

entire cross-section of all the specimens. Micrographs across the entire TGO/bond

coat interfacial area were taken at 5000x magnification for every specimen. High

resolution quantitative elemental analysis was carried out at the TGO/bond coat

interfacial area of all the thermally exposed specimens using EPMA (Electron Probe

Microscope Analysis). A beam current of 20nA was applied with a beam size of about

1m, at an accelerating voltage of 15kV. The scans were done at a step size of 0.1m

in both X and Y directions for an area of 30.00m x 25.00m.

In order to conduct Orientation Imaging Microscopy (OIM) analysis, an additional

polishing step was applied to the specimens using the 0.06 m Buehler Mastermet

colloidal silica suspension to ensure the removal of surface deformation, which is

essential for EBSD analysis. A JEOL FE-SEM equipped with EBSD detector and

OIM data collection software was used to carry out the orientation analysis near the

bond coat/TGO interfacial region of the 0, 30, and 100hrs specimens. The scans were

done at a step size between 0.4 to 0.6 m in both X and Y directions, for an area of

28.000 m x 25.000m. For specimens that had inherently large grains, an additional

scan was conducted using the same step size and area settings as before.

133

Nano-indentation (MTS nano-indentation XP) was carried out along the

cross-sectional surface of bond coat for every specimen. A total of 10 indentions were

conducted at locations approximately 15m below the TGO/bond coat interface on

every specimen, and were separated by a distance of at least 20m from each other.

The indenter load was controlled using a depth limit of 1000nm and a peak holding

times of 10s. The hardness results were used to elucidate the mechanical response of

the bond coat after the cooling stage.

For the purpose of carrying out a more realistic Thermo-Calc analysis of the phase

distribution in the bond coat at the TGO/bond coat interface, energy-dispersive X-ray

spectroscopy (EDX) was utilized to determine the major alloying compositions near

the oxide-bond coat interface (unlike the approximating approach used in the previous

chapter) of the 100hrs isothermally exposed CMSX-4 LT Pt-Al and the 50hrs

isothermally exposed TMS-82+ LT Pt-Al specimens. A total of 23 point scans were

conducted at locations approximately 1-3m below the TGO/bond coat interface on

both specimens in order to avoid collecting secondary electron signals from the TGO

layer (i.e. oxygen and aluminum). Average values of the alloying elements in wt%

were obtained and used in the Thermo-Calc simulation of section 5.3.5. Table 5.3

shows a comparison between the bond coat composition near the TGO/bond coat

interface of the as-coated CMSX-4 LT Pt-Al specimen using the estimation approach

(i.e. Chapter 4) and actual EDX scan (i.e. this chapter).

It can be clearly seen that the estimating approach by proportionally decreasing other

elements in the previous chapter (i.e. dilution effect of aluminization) is flawed, as the

diffusivities tend to vary significantly between different elements. For instance, the W

134

and Ta contents were significantly reduced/diluted in the actual bond coat as

compared to the bulk CMSX-4, while Mo and Ti maintained similar contents to that

of the bulk CMSX-4. In addition, it was assumed that only Al and Ni content varied in

the bond coat due to oxidation, while the rest of the elemental compositions

maintained the same as in their as-coated state. This is definitely not realistic, as

inter-diffusion of certain elements is known to take place during oxidation. It should

be noted; however, that the previous chapter was mainly looking at the relative

amount of Ti solubility in and ’ phases of bond coat. Hence, the discrepancy in the

bond coat compositions should not have a major influence in the general distribution

trend.

Table 5.3

A comparison between the estimated and actual (EDX scanned) bond coat

compositions in as-coated CMSX-4 specimen

Bond coat

composition

Co Cr Mo W Al Ti Ta Re Pt Ni

CMSX-4 LT

Pt-Al

(estimated)

8.31 5.63 0.52 5.54 19.00 0.09 5.69 2.60 N/A Bal.

CMSX-4 LT

Pt-Al

(scanned)

4.96 2.57 0.50 0.93 20.0 0.14 0.76 0.77 30.7 36.1

Substrate

composition

Co Cr Mo W Al Ti Ta Re Pt Ni

CMSX-4 9.6 6.5 0.6 6.4 5.6 1.0 6.5 3 - Bal.

135

5.3 Results

5.3.1 Microscopic observation near the TGO/bond coat interface

Figure 5.2 shows a side by side comparison of the bond coat regions between two 30

and 50hrs isothermally exposed superalloys (SRR99 and CMSX-4) paired with the LT

Pt-Al bond coats. It can be seen that the phase morphology differs significantly

between the two alloys. The SRR99 LT Pt-Al specimen showed short circuit diffusion

along the grain boundaries of grains, and formed a secondary phase of ’

precipitates after 30hrs. In addition, large ’precipitates were also observed near the

TGO/bond coat interfacial region. The corresponding 30hrs isothermally exposed

CMSX-4 LT Pt-Al specimen, on the other hand, maintained a uniform layer of

phase without the presence of ’ precipitates at. After 50hrs of isothermal exposure,

spallation of the TGO had already occurred on the SRR99 LT Pt-Al specimen, while

the TGO layer of the CMSX-4 LT Pt-Al specimen remained intact. Moreover,

CMSX-4 LT Pt-Al was found to contain the martensitic phase (lath structures) along

with phase.

136

Fig. 5.2. SEM micrographs showing a comparison of bond coat regions between 30

and 50hrs isothermally exposed SRR99 LT Pt-Al and CMSX-4 LT Pt-Al specimens,

with a unique ’ formation at the grain boundaries of the SRR99 specimen.

Similarly, Figure 5.3 shows a side by side comparison of the bond coat regions

between 30 and 50hrs isothermally exposed HT Pt-Al and LT Pt-Al bond coats paired

with the same TMS-82+ superalloy. There was a significant difference in the phase

distribution of these two bond coats. The TMS-82+ HT Pt-Al specimen was found to

contain a uniform distribution of martensitic phase (coarse lath structures); while the

TMS-82+ LT Pt-Al specimen maintained mostly phase with minimal formation of

martensites and ’precipitates after 30hrs of isothermal exposure. For the 50hrs

isothermal exposed specimens, large pieces of ’ precipitates began to form on the

TMS-82+ HT Pt-Al specimen beside the martensitic phase. The TMS-82+ LT-Pt-Al,

on the other hand, began to show more distribution of martensites and some ’

precipitates along with the phase.

137

Fig. 5.3. SEM micrographs showing a comparison of bond coat regions between 30

and 50 isothermally exposed TMS-82+ HT Pt-Al and TMS-82+ LT Pt-Al specimens,

with a complete martensitic transformation in the HT-Pt-Al bond coat.

5.3.2 EBSD mapping analyses near the TGO/bond coat interface

EBSD pattern maps are utilized in this section to further highlight the unique features

in bond coat morphology and phase distribution of the TMS-82+ HT Pt-Al and

SRR99 LT Pt-Al specimens, as the EBSD map is able to indicate the grain boundaries

by distinguishing the orientation of individual crystals. It should be noted here that the

EBSD pattern of the TGO layer could not be accurately indexed (most likely due to

its large residual stress induced crystal distortion), and hence, was deliberately

removed from all maps for the ease of viewing. Figure 5.4 shows a side by side

comparison of the grain morphology between the HT Pt-Al and LT Pt-Al specimens

138

paired with the same TMS-82+ superalloy, in the as coated condition. Before

isothermal exposure, the TMS-82+ HT Pt-Al specimen was found to contain much

larger grains than the TMS-82+ LT Pt-Al bond coat. The larger grains can be

explained by the fact that the aluminization processes of the TMS-82+ HT Pt-Al were

carried out at a relatively higher temperature of 1080°C in a low activity aluminum

environment. Based on the fundamental physics of crystallography, grain growth

tends to be much more rapid due to higher diffusivity of atoms at high temperatures.

Fig. 5.4. EBSD mapping showing a comparison of grain morphology of the bond coat

region near the oxide-bond interface of the TMS-82+ HT Pt-Al and TMS-82+ LT

Pt-Al prior to isothermal exposure, with the HT Pt-Al bond coat having much larger

grains.

A similar comparison (Figure 5.5) of 30hrs isothermally exposed specimens shows

that the TMS-82+ HT Pt-Al bond coat transformed completely from the as coated

phase to an acicular martensitic phase, while the TMS-82+ LT Pt-Al specimen

possessed mostly the phase having enlarged grain size. A significant degradation

can be seen in the pattern quality of the 30hrs isothermally exposed TMS-82+ HT

Pt-Al most likely due to crystal distortion due to associated with the significant

139

martensitic phase transformation. Figure 5.6 provides a side by side comparison of the

phase distribution between the SRR99 LT Pt-Al and CMSX-4 LT Pt-Al specimens

after 30hrs isothermal exposure. The SRR99 LT Pt-Al specimen was found to exhibit

the formation of ’ phase at the grain boundary and within the matrix of phase as

secondary precipitate. The CMSX-4 LT Pt-Al map, on the other hand, showed very

uniform grains of clean phase.

Fig. 5.5. EBSD mapping showing the lath martensitic structure of the TMS-82+ HT

Pt-Al specimen, in contrast to the phase of the TMS-82+ LT Pt-Al after 30hrs of

isothermal exposure.

Fig. 5.6. EBSD mapping showing the unique ’ precipitation at the grain boundaries

of the SRR LT Pt-Al, in contrast to the clean and uniform phase of the CMSX-4 LT

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Pt-Al specimen after 30hrs of isothermal exposure.

Figure 5.7, 5.8 and 5.9 show a comparison between the phase distribution of CMSX-4

LT Pt-Al, TMS-82+ LT Pt-Al and TMS-82+ HT Pt-Al after 100hrs of isothermal

exposure. Note that an electron microscopic image is provided to accompany each

individual EBSD map, in order to help identifying the different phases present. The

CMSX-4 LT Pt-Al specimen experienced a transformation from phase to a mostly

coarse martensitic structure along with minor amount of ’phase. The TMS-82+ LT

Pt-Al specimen, on the other hand, showed the presence of more ’phase along with a

coarse martensitic structure. The pattern quality of the martensitic region is

significantly lower relative to the neighboring grains of ’phase. Lastly, the formation

of much larger ’grains could be found in the TMS-82+ HT Pt-Al specimen with

patterns of martensitic regions below recognition.

Fig. 5.7. EBSD mapping showing the degradation in pattern quality of martensitic

region relative to the amount of ’ formation in the bond coat of 100hrs isothermally

exposed CMSX-4 LT Pt-Al.

141

Fig. 5.8. EBSD mapping showing more severe degradation in pattern quality of

martensitic region with increasing amount of ’ formation in the bond coat of 100hrs

isothermally exposed TMS-82+ LT Pt-Al 100hrs.

Fig. 5.9. EBSD mapping showing very severe degradation in pattern quality

associated with even more ’ formation in the bond coat of 100hrs isothermally

exposed TMS-82+ HT Pt-Al.

The EBSD map of the 100hrs isothermally exposed TMS-82+ Pt-diffusion specimen

is compared to that of the 30hrs isothermally exposed SRR99 LT Pt-Al specimen to

highlight the inherent difference between the and ’ structure of the Pt-diffusion

bond coat and the structure of the LT Pt-Al bond coat (see Figure 5.10). It can be

seem that the TMS-82+ Pt-diffusion specimen consisted of small grains of more or

142

less uniform sizes, whereas, the SRR99 LT Pt-Al specimen contained much larger

grains with ’ precipitates of various sizes. Figure 5.11 shows a comparison of the

misorientation profile of the bond coat grains near the TGO/bond coat interface

between the TMS-82+ Pt-diffusion and SRR99 LT Pt-Al specimens. The arrow lines

in Figure 5.10 indicate the location where the line plot was taken from, starting from

the left arrow toward the right one. The flat regions of the plot indicate the location

and width of each individual grain, as the crystal orientation of should be consistent

throughout a grain. The TMS-82+ Pt-diffusion specimen consists of grains having

consistent size and orientations, while sudden fluctuation in the orientation can be

seen within grain filled with many precipitates fine grains in the SRR99 LT Pt-Al

specimen.

Fig. 5.10. EBSD mapping highlighting the inherent difference between the ’

microstrucuture of the Pt-diffusion bond coat and the microstructure of the LT Pt-Al

bond coat.

143

Fig. 5.11. A comparison of the misorientation profile near the oxide-bond coat region

between the TMS-82+ Pt diffusion (top) and SRR99 LT Pt-Al (bottom) after 100hrs

and 30hrs of isothermal exposure, with much sharper peaks in the SRR99 specimen

144

indicating regions of ’ precipitation.

5.3.3 EPMA mapping analysis near TGO/bond coat regions

EPMA analysis was carried out to study the 30hrs and 100hrs isothermally exposed

TMS-82+ Pt-diffusion, in order to study the chemical diffusion and segregation

behavior of its bond coat (see Figure 5.12). After 30hrs of isothermal exposure, the Al

content near the TGO/bond coat region of the TMS-82+ Pt-diffusion specimen was

uniformly depleted with corresponding enrichment in the Cr content. The Ti content

was found to segregate below the Cr enriched layer at the exact location where Cr

depletion occurred. After 100hrs of isothermal exposure, the Ti and Cr chemistry

remained very similar to that of the 30hrs isothermally exposed while further Al

depletion took place. It should be noted that traces of chromium oxide can be found in

the TGO layer for both exposure times.

145

Fig. 5.12. EPMA mapping showing the chemistry near the oxide-bond coat regions of

the TMS-82+ Pt-diffusion specimens after 30 and 100hrs of isothermal exposure, with

uniformly depleted layer of Al and Ti while enriched with Cr.

Figure 5.13 provides a direct comparison of the Ti and Cr distribution between the

TMS-82+ LT Pt-Al and the CMSX-4 LT Pt-Al specimens having isothermally

exposed for 100hrs. Similarly, Ti-Cr segregation was found in the TMS-82+ LT Pt-Al

specimen, except that the segregation occurred in small patches along bond coat

region near the TGO/bond coat interface rather than layers as in the previous case. On

the other hand, the CMSX-4 LT Pt-Al specimen showed a rather consistent Cr

distribution with much less Ti content near the interface. Chromium oxide was

detected in the TGO layers of both specimens.

146

Fig. 5.13. EPMA mapping showing a comparison of Cr content near oxide-bond coat

interface between TMS-82+ LT Pt-Al and CMSX-4 LT Pt-Al after 100hrs of

isothermal exposure, with CMSX-4 showing no obvious Ti presence.

5.3.4 Hardness by nano-indentation near the TGO/bond coat interface

To study the mechanical property of the bond coat region near the TGO/bond coat

interface, hardness results measured by nano-indentation is plotted in Figure 5.14 here

for the 30, 50 and 100hrs isothermally exposed specimens. Note the error bar here

represents the standard deviation of 10 measurements for each specimen. After 30hrs

of isothermal exposure, the SRR99 LT Pt-Al specimen had the highest hardness and

standard deviation. In general, all specimens except the TMS-82+ Pt-diffusion,

underwent an increasing trend in both average hardness and standard deviation

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between 50 and 100hrs. The TMS-82+ HT Pt-Al had the highest hardness and

standard deviation among all specimens listed here after 100hrs of isothermal

exposure. The TMS-82+ Pt-diffusion specimen was found to maintain similar

hardness values and standard deviation going from 30 to 100hrs.

Fig. 5.14. A plot showing the hardness values by nano-indentation of various bond

coats near the oxide-bond coat interface, with increasing trend in both hardness values

and standard deviations.

5.3.5 Simulation of phase transformation by Thermo-Calc

To clarify the effect of Cr content on the phase distribution of bond coat, a

thermodynamic software, named Thermo-calc was used to identify the phase

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distribution of the LT Pt-Al bond coat system when prepared on the CMSX-4 alloy. A

set of parameters including the isothermal temperature used (1409K), and the bond

coat compositions near the TGO/bond coat interface of the 100hrs isothermally

exposed CMSX-4 LT Pt-Al specimen were specified in the calculation. We assume

that the system is in a stable equilibrium state after certain amount of time. Note that

the bond coat compositions were taken from the EDX analysis result as mentioned in

the experimental section (see Table 5.4).

Table 5.4

The nominal composition in wt% of the CMSX-4 LT Pt-Al and TMS-82+ LT Pt-Al

bond coats as specified in the Thermo-Calc calculation

Bond coat

composition

Co Cr Mo W Al Ti Ta Re Pt Ni

CMSX-4 LT

Pt-Al

6.691 3.158 0.502 1.998 11.117 0.0104 5.345 0 20.429 50.211

TMS-82+

LT Pt-Al

4.898 2.887 0.476 1.863 11.787 0.316 3.704 0 22.531 50.902

Figure 5.15a demonstrates the mole-fraction of different phases near the TGO/bond

coat interfaces of the 100hrs isothermally exposed CMSX-4 LT Pt-Al alloy based on

the calculation in Thermo-Calc as a function of Cr content. The results here showed

an increase in the mole fraction of phase corresponding to a decrease in the ’ with

increasing Cr content. The dotted line shows the relative distribution of the bond coat

phases in the 100hrs isothermally exposed CMSX-4 LT Pt-Al specimen. A similar

analysis was carried out to study the effect of Ta content on the phase distribution of

the 50hrs isothermally exposed TMS-82+ LT Pt-Al. The purpose here is to see if a

low Ta content, such as in the case of SRR99 LT Pt-Al specimen, could have an

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influence on the ’ transformation behavior. Again, the bond coat compositions used

for Thermo-Calc analysis were taken from EDX line scans near the oxide-bond coat

interface (see Table 5.4). Thermo-calc results given in Fig. 5.15b also indicated an

increase in the mole fraction of phase corresponding to a decrease in the ’ with

increasing Ta content. However, the extent of the phase variation seemed much less in

this case.

Fig. 5.15. The mole fraction of different phases existing near the TGO/bond coat

interface of the bond coats in a) CMSX-4 LT Pt-Al 100hrs, and b) TMS-82+ LT Pt-Al

50hrs specimens as a function of varying Cr and Ta contents, respectively. These

elements generally help stabilizing phase.

5.4 Discussion

5.4.1 Interfacial evolution during thermal exposure

Table 5.5 is a summary of the phase evolution of the five TBC systems with

corresponding hardness of bond coats between 30 and 100hrs isothermal exposures,

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and their TGO stress obtained from the previous chapter. The phase constituents of

the bond coat seem to have a huge influence on the TGO stress level. The TMS-82+

Pt-diffusion specimen, having pure phase between 30 and 100hrs isothermal

exposures, was measured to have the highest average hardness values, while having

the highest TGO stress throughout thermal exposure. The CMSX-4 LT Pt-Al

specimen with pure phase after 30hrs of isothermal exposure had a slightly lower

TGO stress than that of the TMS-82+ HT Pt-Al with all martensitic structure at 30hrs

(~2.29GPa vs. 2.45GPa).

The previous chapter indicated that the CMSX-4 LT Pt-Al and TMS-82+ HT Pt-Al

were the only specimens to have experienced a decline in TGO stress between 30 and

80hrs of isothermal exposure, and in the case of CMSX-4 LT Pt-Al, the decline

continued all the way to 100hrs. The TMS-82 LT Pt-Al, on the other hand, started out

with relatively lower TGO stress having multiple phases after 30hrs isothermal

exposure, and continued to show an increasing trend in TGO stress until 100hrs of

isothermal exposure. These phenomena suggests that a transformation from pure to

multi-phases can largely alleviate the TGO stress in the systems, despite having

high stress inducing phases such as martensites and ’.

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Table 5.5

A summary of the phase evolution, hardness, and TGO stress of all the bond coats

studied between 30 and 100hrs of isothermal exposure

Phase evolution, hardness, and TGO stress during thermal

exposure

30hrs 50hrs 100hrs

SRR99 LT Pt-Al ’ precipitates (g.b.

and near interface)

+

Stress : 0.237GPa

Hardness:

7.197GPa

-N/A (Spalled)

-N/A

TMS-82+ HT

Pt-Al

All lath martensitic

structure

Stress: 2.446GPa

Hardness:

4.555GPa

Martensitic + ’

(significant)

Stress: 2.327GPa

Hardness:

4.825GPa

Martensitic + ’

(large amount)

Stress: 2.485GPa

Hardness:

8.290GPa

TMS-82+ LT

Pt-Al

Mostly +

martensites + ’

(trace)

Stress: 1.183GPa

Hardness:

4.520GPa

martensites

(significant) + ’

(more trace)

Stress: 1.775GPa

Hardness:

4.679GPa

Martensitic +’

(significant)

Stress: 1.854GPa

Hardness:

5.322GPa

CMSX-4 LT Pt-Al All

Stress: GPa

Hardness:

5.167GPa

martensites

(significant)

Stress: 2.170GPa

Hardness:

4.564GPa

Martensitic + ’

(significant)

Stress: 2.091GPa

Hardness:

5.733GPa

TMS-82+

Pt-diffusion

Uniform layer

near the interface

Stress: 2.682GPa

Hardness:

6.553GPa

Uniform layer

near the interface

Stress: 2.840GPa

Hardness:

6.315GPa

Uniform layer

near the interface

Stress: 2.880GPa

Hardness:

6.41GPa

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5.4.2. Phase transformation due to substrate/bond coat chemistry

In TBCs, Al diffuses faster through the grain boundaries of the bond coat upon

forming Al2O3 during thermal exposure. However, the question still remains as to why

the formation of the ’ phase along the grain boundaries was exclusively observed on

the SRR99 LT Pt-Al in early isothermal times. As previously confirmed by EDX,

EPMA, and Thermo-Calc simulation, Ti and Ni contents have the tendency to

segregate to Al depleted region, corresponding to the location of ’ phase. The SRR99

LT Pt-Al specimen, having a relatively higher Ti and Ni content (see Table 5.1), was

previously measured to have the fastest TGO growth kinetic and the ’ phase

formation rate. Both Ni and Ti sped up the Al depletion process through the grain

boundaries of the SRR99 bond coat, thereby, migrated to the grain boundary to form

the ’ phase.

The unique phase transformation phenomena observed on the SRR99 LT Pt-Al

specimen is believed to be the cause of its relatively poor interfacial adhesion and

spallation lifetime. EPMA results in Figures 5.1 and 5.13 of the SRR99 LT Pt-Al and

TMS-82+ LT Pt-Al specimens, respectively, confirmed traces of Ti in the TGO layer,

while Ti rich ’ precipitates were present near the TGO/bond coat interface. The

detrimental effect of TiO2 formation near the bond coat surface, and subsequently,

degrading the adherence of the alumina scale is well documented [18,113,114]. The

same degradation mechanism, however, does not seem to apply in the case of the

TMS-82+ HT Pt-Al specimen. Despite having Ti containing ’ precipitates near the

TGO/bond coat interface, no obvious Ti was detected (see Figure 5.1) in the TGO

scale as confirmed at multiple locations along the interface. The amount of Ti in the ’

precipitates of the TMS-82+ HT Pt-Al was considerably lower than that in the SRR99

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LT Pt-Al specimen, which may be the reason behind the absence of TiO2.

Thermo-Calc simulation in section 5.3.5 indicated that Cr and Ta tend to stabilize the

phase, preventing the early formation of ’. However, the effect of Cr seemed to be

much greater than that of Ta. A comparison of the EPMA compositional mapping

between the TMS-82+ LT Pt-Al and CMSX-4 LT Pt-Al specimens (Figure 5.13)

revealed that Cr, if having sufficient concentration, can largely inhibit the diffusion of

Ti content outward to the bond coat surface. In fact, the CMSX-4 LT Pt-Al, having

higher Cr content in the bulk alloy (see Table 5.1), was shown in the previous chapter

to have relatively slower ’ formation despite having more bulk Ti content than the

TMS-82+ LT Pt-Al specimen. It should be noted that the stability of a certain phase

may as well be influenced by other alloying elements in the bulk superalloy.

Nonetheless, the results suggested that Cr plays a critical role in the prevention of

early ’ formation.

The complete phase transformation to the coarse martensitic structure in the HT Pt-Al

bond coat is believed to be associated with its martensitic transformation temperature,

Ms. A previous work by Smialek et al [43] indicated that Ms in -NiAl alloy increases

when Al depletion occurs as a result of oxidation and alloy/bond coat inter-diffusion.

As pointed out by Hangen et al [115], the martensitic transformation does not undergo

complete (i.e. 100%) transformation right below the Ms, but instead continues to

increase in volume fraction with decreasing temperature (i.e. increasing cooling). In

the previous EDX line profile study of the Pt-Al bond coats by Wu et al [35], the Al

content in the HT Pt-Al bond coat within a distance of 20m from the TGO/bond coat

interface was determined to be slightly lower (~7 at % lower on average) than that in

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the LT Pt-Al. These findings suggest that the complete transformation to coarse

martensite in the HT Pt-Al bond coat type is most likely due to its inherently low Al

content, hence, having a higher Ms.

5.4.3. Phase transformation induced interfacial rumpling

Qualitative and semi-quantitative assessments of local plastic strain by utilizing both

the changes in the EBSD pattern quality and grain misorientations have long been

conducted in many previous studies [116-118]. It is believed that the progressive

degradation in the pattern quality of the martensitic regions corresponds to an

increasing severity of plastic strain. The HT XRD results in the work by Chen et al

[44] indicated that the molar volume of the L10 martensite is approximately 2%

smaller than that of the , causing a transformation strain of about -0.7% during

cooling. They, along with several other researchers [119], believe that this volume

shrinkage due to martensitic transformation can generate additional stress during

thermal cycling in the bond coat, and thus, can significantly enhance rumpling

growth.

The role of martensitic transformation in causing rumpling should be questioned in

the case here with isothermal single cool-down tests. As pointed out in previous

literatures [39], rumpling amplitude tends to accumulate due to an intrinsic

asymmetry in the plastic response to tension and compression of the bond coat (i.e.

during cooling and heating). Rumpling tends to accumulate slowly, as the amount of

undulation developed per cycle is rather small. It should be noted; however, that the

155

martensitic transformation may contribute to lowering interfacial fracture energy by

creating out-of-plane stresses at the bond coat/TGO interface (i.e. crack opening).

Although, the volume shrinkage associated with to ’ phase transformation is

significantly larger than that of to martensitic phase (i.e. 8 to 38% change depending

on whether Al depletion occurred predominantly by oxidation or inter-diffusion) [39],

it should be noted; however, that the process takes place at a much higher temperature

range, in which the creep rate is high. As a result, much of the TGO growth stress is

relaxed in the form of rumpling, as evident from the low stress level and high

rumpling rate of TMS-82+ LT Pt-Al specimen shown in the last chapter. In contrast,

the to martensitic transformation occurred in a significantly shorter time span (i.e.

cooling) at a much lower temperature range (i.e. less creep relaxation), in which the

strain generated by the martensitic transformation together with the CTE mismatch

between the coating layers contributes as additional elastic strain at the bond

coat/TGO interface without inducing rumpling. Hence, it is believed that the effect of

to martensitic transformation is more significant in triggering crack nucleation at the

bond coat/TGO interface.

EBSD results of 100hrs isothermally exposed specimens revealed that the degradation

of the martensites patterns tend to be the most severe when large ’ precipitates were

formed as neighboring grains (see Figure 5.7-5.9). Based on this observation, it is

proposed that the adjacent ’grains, along with the TGO scale, may impede the

contraction of martensites during cooling by imposing a tensile strain. This force

exerted on the martensites is likely to be proportional to the number of ’ precipitates

formed in the surroundings. Hence, TMS-82+ HT Pt-Al, having the fastest ’

formation at 100hrs into isothermal exposure, was shown to have the most severe

156

degraded pattern in the martensites. On the other hand, the CMSX-4 LT Pt-Al

specimen with significantly less ’ formation than the TMS-82+ LT PT-Al after

100hrs of isothermal exposure, had noticeably less degradation in its EBSD map.

Note that the amount of ’ precipitates formation near the TGO/bond coat interface

was thoroughly quantified on these TBC systems in the previous chapter, and used as

reference for comparison here.

Tolpygo et al argued in their previous work [39] that reverse martensitic

transformation during cyclic thermal exposure has no discernible effect on the

rumpling of bond coat. This argument, however, does not conform to our results

shown in this work. As discussed earlier, when martensites were found as the

dominant phase (volume %) in the bond coat, such as in the case of the 30hrs

isothermally exposed TMS-82+ HT Pt-Al and the 100hrs isothermally exposed

CMSX-4 LT Pt-Al 100hrs specimens, there was much less pattern degradation (i.e.

severity of plastic strain) in their corresponding EBSD mappings as oppose to the

cases where ’ phases co-existed with martensites (i.e. TMS-82+ HT Pt-Al and

TMS-82+ LT Pt-Al at 100hrs). Moreover, their assumption that no martensitic

transformation occurred in their two-zone cycling temperature range lacks the

consideration that the Ms temperature increases with Al depletion and interdiffusion

of various alloying additions during oxidation, hence, should be considered as a

dynamic parameter [115,116,120]. Clearly, the additional strain generated by the

martensitic transformation may cause rumpling with increasing thermal cycles when

the Ms temperature is high enough for creep relaxation to take place in the bond coat.

The TMS-82+ Pt-diffusion specimen was shown to undergo a uniform Al depleted

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layer near the TGO/bond coat interface, with corresponding enrichment in Cr. The

segregation of Ti was also observed right beneath this Al-depleted and Cr-enriched

layer. This compositional distribution profile seems to be maintained between 30hrs

and 100hrs of isothermal exposure, while the as-coated ’ + phase gradually

transformed to a single phase microstructure. In a previous work [35], the TMS-82+

Pt-diffusion specimen was found to exhibit no rumpling behavior throughout its entire

thermal cycling test, in comparison with other TBC systems. In addition, the

TMS-82+ Pt diffusion specimen had some of the best interfacial adhesion among the

TBCs as indicated in the last chapter, despite having the largest TGO stress of the

entire specimen set during thermal exposure. The reason behind these phenomena can

be related to the uniform layer of phase microstructure near the bond coat/TGO

interface. Unlike the to ’/martensitic transformation of Pt-Al bond coat systems,

there is no significant volume change associated with ’ to transformation, and

hence, does not produce a rumpling growth. Moreover, the Pt-diffusion systems do

not have Ti segregation near the bond coat/TGO interface, and thus would not form

TiO2 in the TGO layer. Contrary to the Pt-diffusion system, the TMS-82+ LT Pt-Al

and TMS-82+ HT Pt-Al systems had relatively poor interfacial adhesion/spallation

lifetime and continuous growing rumpling amplitude during thermal exposure (as

shown in last chapter).

It is proposed here that the co-existence of martensite, and ’ phases with inherently

very different thermomechanical properties may induce a misfit stress in the TGO

layer during thermal cycling, leading to crack formation at the TGO/bond coat

interface and thus coating spallation. This idea is supported by the increasing trend of

the hardness and standard deviation values measured from various bond coat systems

158

shown in section 5.3.4. Though not provided in the plot before, the average hardness

of as-coated bond coat systems (i.e. as coated CMSX-4 LT Pt-Al, TMS-82+ HT

Pt-Al and LT Pt-Al) was measured to be 5.42GPa. When the bond coat transformed to

mostly martensitic phase (i.e. TMS-82+ HT Pt-Al after 30hrs of isothermal exposure,

TMS-82+ LT Pt-Al and CMSX-4 LT Pt-Al after 50hrs of isothermal exposure), the

average hardness became 4.60GPa. As more and more ’ phase began to precipitate

out alongside the martensitic structure after 100hrs of isothermal exposure, a drastic

increase in both hardness and standard deviation values from those of the 50hrs

thermally exposed samples can be seen for all Pt-Al systems, especially for the

TMS-82+ HT Pt-Al specimen (see Figure 5.14). These variations in the hardness

value due to different phases are likely to induce a misfit stress in the TGO layer

across the TGO/bond coat interface, as the resistance to plastic deformation during

thermal cycling differs locally in the bond coat.

Specimens that partially underwent to ’ phase transformation were found to have

the highest degrees of rumpling. For instance, the SRR99 LT Pt-Al, TMS-82+ LT

Pt-Al and CMSX-4 LT Pt-Al specimens, all of which underwent partial to ’ phase

transformation, had the highest rumpling growth rate among all specimens between

as-coated state and 100hrs isothermal exposure. In contrast, the TMS-82+ Pt-diffusion

and TMS-82+ HT Pt-Al, both of which underwent a complete phase transformation at

early isothermal hours, had relatively lower rumpling growth rate (See previous

chapter).

159

5.4.4 Driving force for spallation versus interfacial fracture toughness

The results of this work have clarified the difference in the degradation behavior

between the SRR99 LT Pt-Al and TMS-82+ HT Pt-Al specimens after 30 and 100hrs

of isothermal exposure, respectively. Figure 5.16 provides a schematic explanation of

the TGO stress evolution and the interfacial degradation process during isothermal

heating and subsequent cooling of the two specimens. During isothermal holding,

TGO began to thicken on both specimens; however, the SRR99 LT Pt-Al had a much

faster oxidation kinetic with traces of TiO2 present at the TGO/bond coat interface

while the TMS-82+ HT Pt-Al specimen maintained mostly pure alumina. The lateral

TGO growth due to the formation of new oxide within the scale interior was

constrained by the underlying metal substrate, which led to an in-plane compression

of the oxide scale. This existence of lateral growth strain at isothermal temperature is

known to cause creep in both the scale and bond coat, and thus induces surface

undulation [40]. In addition, the volume shrinkage associated with ’ formation led to

a local stress concentrated zone in the TGO above the ’ grain. This uneven strain

response between the and ’ regions would further enhance rumpling growth.

160

Fig. 5.16. A qualitative illustration of the TGO stresses evolution and interfacial

degradation during isothermal heating and subsequent cooling of the 30hrs

isothermally exposed SRR99LT Pt-Al and the 100hrs isothermally exposed TMS-82+

HT Pt-Al. The SRR99 had more rumpling due to ’ formation at the grain boundary,

while the TMS-82+ had a relatively higher residual stress in the TGO due to

martensitic transformation.

It is interesting to note that despite having a thicker TGO and more ’ formation, the

161

TMS-82+ HT Pt-Al after 100hrs isothermal oxidation had significantly less rumpling

compared to the SRR99 LT Pt-Al after 30hrs isothermal oxidation (See last chapter).

As pointed out in a previous work by Tolpygo [38], the bond coat was observed to

swell during high temperature oxidation due to Kirkendal effect, leading to surface

undulation of the bond coat. There seemed to be uneven response to bond coat swell,

where the inner grains of bond coat swelled more (i.e. undulation peaks) than the

grain boundaries (i.e. undulation valleys). Based on this phenomenon, it is believed

that unique ’ formation along the grain boundaries of the SRR99 specimen further

contributed to additional surface undulation due to the volume shrinkage (i.e. uneven

strain) of to ’ transformation, making the specimen more rumpled.

During cooling, TGO stress began to increase due to the CTE mismatch between the

different phases of the bond coat and the TGO layer. The bond coat contracted more

and thus applied a compressive stress on the TGO, while the TGO induced a tensile

stress on the bond coat. The TMS-82+ HT Pt-Al was different from the SRR99 LT

Pt-Al in that its phase began to transform into a coarse martensitic phase,

accompanied by molar volume shrinkage during cooling. This transformation likely

led to a higher stress-build up in the TGO layer above the martensitic phases relative

to the regions above ’ in the HT Pt-Al specimen at room temperature, which may

cause crack opening at the TGO/bond coat interface. This is evident from the

relatively higher TGO stress of the 30hrs isothermally exposed TMS-82+ HT-Pt-Al,

having a poor interfacial adhesion as determined previously. The SRR99 LT Pt-Al, on

the other hand, had much less stress gradient across the TGO, where regions above

the had more stress relaxation than those above the ’, as was more susceptible to

creep at high temperature. The main reason for the early spallation of the SRR99 LT

162

Pt-Al can be associated with its faster rumpling growth, while the formation of the

less adhesive TiO2 at the TGO/bond coat interface likely contributed to the spallation.

As shown in the previous chapter, while SRR99 LT Pt-Al attained the highest

rumpling magnitude and ’ formation rate prior to spallation, other specimens (i.e.

TMS-82+ LT Pt-Al and CMSX-4 LT Pt-Al after 100 and 200hrs of isothermal

exposure) were found to gain even higher rumpling magnitude and more ’ formation,

yet continued to survive well into 200hrs of isothermal exposure.

In the previous chapter, a semi-quantitative plot of the variation of the driving force G

for spallation and the interfacial fracture toughness Gc with isothermal exposure was

given to show the remaining lifetime of the TBCs used in this work between 30 and

100hrs. The intention of that plot was to highlight the concept that the interfacial

adhesion of TBCs can only be determined by taking into account the dynamic change

of G and Gc with time and temperature. This concept is again demonstrated by the

degradation behavior of the SRR99 LT Pt-Al and TMS-82+ HT Pt-Al specimens

shown in this work. For instance, the formation of TiO2 at the TGO/bond coat

interface of the SRR99 LT Pt-Al, resulting in the loss of interfacial adhesion, can be

treated as a factor for driving the Gc value lower during isothermal exposure. On the

other hand, the martensitic transformation of the TMS-82+ HT Pt-Al during cooling,

leading to a stress build-up in the TGO, may be considered as a factor behind an

increasing G value.

163

5.5 Conclusion

The results of this work have led to the following conclusions:

The interfacial degradation behavior has been associated with the phase

transformation of bond coat during thermal cycling. This is reflected by the

variation of interfacial rumpling and TGO stress with bond coat hardness

values, which are dependent on the phases present.

The progressive degradation in the EBSD pattern quality of the martensitic

regions due to increasing ’ phases nearby has suggested that a misfit stress

may be generated in the TGO layer due to the inherently different

thermomechanical properties between different phases.

The experimental results have clarified the difference in the degradation

mechanism between the SRR99 LT Pt-Al and TMS-82+ HT Pt-Al specimens.

Early spallation of the SRR99 LT Pt-Al was most likely caused by its fast

rumpling rate due to ’ formation at the grain boundary, while formation of

weakly adherent TiO2 scale due to Ti-rich ’ precipitates near the bond

coat/TGO interface likely contributed to the interfacial degradation. The early

spallation of the TMS-82+ HT Pt-Al was mainly driven by the stress build-up

in the TGO as a result of the volume shrinkage (i.e. out-of-plane stress)

associated with the martensitic transformation.

Thermo-Calc and EPMA analysis indicated that Cr and Ta, if given high enough

amount, has the effect of preventing diffusion of Ti toward to the TGO/bond

coat interface by stabilizing the and/or phases.

164

CHAPTER 6

Summary and suggested future work

6.1 Summary

Current numerical approaches in modeling the intrinsic failure of TBC relies largely

on the notion that spallation occurs when the accumulating strain energy stored in the

coating exceeds a fixed critical value resembling the interfacial adhesion. If this is to

be entirely correct, one would expect this critical value of interfacial adhesion varies

with different materials, but stays independent of their thermal exposure history. In

this study, a unique cross-sectional indentation technique was developed to

quantitatively characterize the adhesion of oxide-bond coat interface among 5

systematically prepared material systems. The results not only re-confirmed that

interfacial adhesion is a material specific property in general, but more importantly,

conclusively demonstrated that the adhesion is dynamic, in particular with time and

temperature. With an aim of further understanding the dynamics (i.e. establishing

correlations between time and temperature dependent microstructure effects),

parameters such as the oxide growth rate, rumpling of the oxide-bond coat interface,

and phase transformation of bond coat were studied as a function of thermal exposure.

It has been conclusively demonstrated here that the oxide-bond coat interfacial

adhesion bears strong dependency on the phase distribution of the bond coats and

TGO growth rate, while receiving little influence from TGO rumpling and residual

stress.

165

The influence of substrate and bond coat chemistry on the degradation mechanisms

leading to the early spallation of thermal barrier coatings (TBCs) has not been well

understood despite years of research effort. This is largely due to the sheer number of

factors (i.e. interfacial rumpling, thermally grown oxide (TGO) growth kinetics, and

TGO residual stress, etc.) that all seem to contribute to the degradation of TBCs. To

clarify the chemical effect, a complete investigation, utilizing the EBSD pattern

variation of bond coat near the oxide-bond coat interface, was carried out along with

various other characterizations on the isothermally exposed TBC specimens.

Specimens that underwent partial phase transformation from to’ near the interface

were determined to have larger rumpling magnitude with lower TGO stress, as

compared to those that transformed completely to a single phase (i.e. Pt-diffusion

type). It is evident that the formation of ’ along the grain boundaries can significantly

enhance rumpling, while martensitic transformation during cooling creates

out-of-plane stresses at the oxide-bond coat interface. These two degradation

mechanisms are likely to be the main reason behind early spallation of TBCs. To

alleviate these two mechanisms, it is necessary to minimize the formation of ’ at the

oxide-bond coat interface. Cr and Ta contents in as-fabricated substrate were

determined to stabilize and phases (i.e. inhibiting fast γ’ formation). Having less ’

formation near the interface can also prevent the formation of weakly adherent

titanium oxides, since Ti preferentially segregates to the γ’ phase.

166

6.2 Suggested future work

By examining the results obtained from the works of Chapter 4 and 5, it can be

summarized that the spallation lifetime of TBCs is dictated by a work of many

different degradation mechanisms (i.e. TGO growth kinetics, type of TGO species,

and martensitic transformation, etc.). The key to understand these factors lies in the

backward trace of key elemental influence from bond coat and substrate chemistry.

Despite many significant findings on the degradation mechanism of TBCs in this

thesis, many aspects of the results can be further refined by carrying out the following

list of work.

- Fabricate Ti, Cr, and Ta modified nickel aluminide coatings (i.e. ternary Ni-Al-X

or quarternary Ni-Al-Pt-X alloys, where X indicates minor elements such as Ti, Cr

and Ta). Similar microstructural studies could be carried out on these types of

specimens after subjecting them to thermal cycling for different length of times.

XRD (X-ray diffraction) analysis can be utilized to identify the overall oxide

compositions formed on the surface. In addition, EPMA can be used to determine

the differences in local chemistry between the different layers and regions of

oxide formation. It will be particularly interesting to identify and compare the

oxide species grown above different regions of bond coats, where the existing

phases are different from each other (i.e. regions of , ’, and martensites, etc.).

EBSD analysis can be applied to identify the different phases of the bond coat

near the oxide-bond coat regions.

- Try to characterize the oxide-bond coat interface of TBCs. with the high

resolution capability of TEM. It would be interesting to study the oxide phases

167

and crystal orientations above different regions of bond coats having different

phases (i.e. regions of , ’, and martensites, etc.). Specimens containing

oxide-bond coat interfaces of different regions can be prepared using FIB-INLO

technique (Focused Ion Beam In-Situ Lift-Out).

- Despite the success of a semi-quantitative approach to ranking the interfacial

adhesion of TBCs in this work, the results are limited in that a cross comparison

between coatings of other material systems is not possible. As mentioned in the

literature review chapter, a universal damage parameter on the basis of interfacial

fracture toughness is necessary to quantify the interfacial adhesion between any

two materials. It would be worthwhile to further refine the approximation of

interfacial fracture toughness, Gc, in section 4.4.1, so that a more accurate

toughness results can be obtained.

168

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