study of bond coats and failure mechanisms of thermal
TRANSCRIPT
STUDY OF BOND COATS AND FAILURE
MECHANISMS OF THERMAL BARRIER
COATINGS
A thesis submitted to the University of Manchester for the degree of
Doctor of Philosophy
in the Faculty of Science and Engineering
2019
Chen LIU School of Materials
LIST OF CONTENTS
1
List of Contents
List of Contents .................................................................................................................... 1
List of Figures ....................................................................................................................... 5
List of Tables .......................................................................................................................11
List of Abbreviations.............................................................................................................12
List of Publication ................................................................................................................13
Abstract .............................................................................................................................14
Declaration .........................................................................................................................15
Copyright Statement ............................................................................................................16
Acknowledgements..............................................................................................................17
Chapter 1 Introduction .........................................................................................................18
1.1 Gas-turbine engine materials ........................................................................................18
1.2 Introduction of thermal barrier coatings .........................................................................19
1.3 Objectives and structure of the dissertation ....................................................................22
Chapter 2 Literature Review ..................................................................................................25
2.1 Thermal barrier coating system .....................................................................................25
2.2 Top coat ....................................................................................................................25
2.2.1 Material requirement and selection.......................................................................25
2.2.2 Deposition techniques and microstructure .............................................................29
2.2.3 New top coat candidates .......................................................................................31
2.3 Bond coat ..................................................................................................................32
2.3.1 Material requirements .........................................................................................32
2.3.2 Bond coat categories ............................................................................................33
2.3.2.1 β-NiAl based diffusion bond coat .......................................................................33
2.3.2.2 Pt-diffused γ-Ni/γ’-Ni3Al bond coat ....................................................................38
2.3.2.3 MCrAlY overlay bond coat ................................................................................41
2.4 Thermally grown oxide.................................................................................................43
2.4.1 Material requirements and selection .....................................................................43
2.4.2 TGO microstructure and stress ............................................................................44
2.4.3 TGO transformation during early stage of oxidation ..............................................46
2.5 Superalloy substrate ....................................................................................................50
LIST OF CONTENTS
2
2.5.1 Composition and microstructure ..........................................................................51
2.5.2 Physical and mechanical properties ......................................................................52
2.6 The degradation and failure of TBCs...............................................................................53
2.6.1 General failure modes..........................................................................................55
2.6.2 The failure mechanisms of APS TBCs ...................................................................58
2.6.3 The failure mechanisms of EB-PVD TBCs.............................................................60
2.6.4 Interface toughness measurement of TBCs ............................................................64
2.6.4.1 Definitions of interface adhesion & delamination and interface toughness ..............64
2.6.4.2 Interface toughness test methods .....................................................................67
2.7 Summary ...................................................................................................................70
Chapter 3 Pt Effect on Early Stage Oxidation Behaviour of Pt-diffused γ-Ni/γ’-Ni3Al Coatings ..........71
3.1 Introduction ...............................................................................................................71
3.2 Experimental procedures .............................................................................................73
3.2.1 Sample preparation and thermal treatment ...........................................................73
3.2.2 Luminescence measurement and data processing ...................................................75
3.2.3 ASTAR automated crystal orientation mapping on TEM .......................................76
3.2.4 Other characterization methods ...........................................................................77
3.3 Results ......................................................................................................................78
3.3.1 Microstructure of the as-fabricated coatings with different Pt contents ...................78
3.3.2 Transient alumina to stable α-Al2O3 transformation...............................................79
3.3.3 TGO composition and microstructure evolution ....................................................82
3.3.4 TGO growth rate & stress evolution .....................................................................84
3.3.6 PLPS studies on Ni-Al-Pt alloy samples .................................................................89
3.4 Discussion ..................................................................................................................94
3.4.1 Pt effect on the θ-Al2O3 to α-Al2O3 transformation .................................................94
3.4.2 Pt effect on TGO composition & stress ..................................................................98
3.4.3 Early stage oxidation effect on prolonged oxidation performance ............................99
3.5 Summary ................................................................................................................. 100
Appendix A. Coating/alumina orientation analysis............................................................... 100
Appendix B. Prolonged oxidation lifetime of coatings with different Pt additions ..................... 104
Appendix C. Early stage oxidation effect on the stable scale morphology & stress .................... 105
Chapter 4 Effect of Superalloy Substrate on the Lifetime and Interfacial Toughness of Electron Beam
Physical Vapour Deposited Thermal Barrier Coatings .............................................................. 107
4.1 Introduction ............................................................................................................. 107
LIST OF CONTENTS
3
4.2 Experimental procedures ........................................................................................... 109
4.2.1 Sample preparation ........................................................................................... 109
4.2.2 Thermal treatment ............................................................................................ 110
4.2.3 Microstructure characterization ......................................................................... 110
4.2.4 Interface toughness measurement by the strain-to-fail test.................................... 111
4.2.4.1 Theoretical background ................................................................................. 111
4.2.4.2 Strain-to-fail test coupled with 3D-DIC ............................................................. 112
4.2.4.3 Determination of the coating stress ................................................................. 113
4.2.4.4 Determination of the YSZ modulus .................................................................. 114
4.2.4.5 Measuring buckling radius by 3D-DIC ............................................................... 114
4.3 Results .................................................................................................................... 115
4.3.1 Microstructure of the as-received TBCs .............................................................. 115
4.3.2 Cyclic oxidation testing ...................................................................................... 116
4.3.2.1 YSZ lifetime .................................................................................................. 116
4.3.2.2 Microstructural evolution ............................................................................... 117
4.3.3 Strain-to-fail compression test coupled with 3D-DIC ............................................ 123
4.4 Discussion ................................................................................................................ 127
4.4.1 Estimation of the interfacial toughness for TBCs on N5 and X4 substrates............. 127
4.4.2 Interface degradation of TBCs on different substrates.......................................... 129
4.5 Summary ................................................................................................................. 132
Chapter 5 The Al-enriched γ’-Ni3Al-base bond coat for thermal barrier coating applications ......... 133
5.1 Introduction ............................................................................................................. 133
5.2 Experimental procedures ........................................................................................... 136
5.2.1 Sample preparation ........................................................................................... 136
5.2.1.1 Fabrication of Pt-diffused intermediate coatings................................................ 136
5.2.1.2 Fabrication of Al-enriched γ’-phase coatings by pack cementation ....................... 138
5.2.2 Thermal treatment ............................................................................................ 140
5.2.3 Characterization methods .................................................................................. 140
5.3 Results .................................................................................................................... 141
5.3.1 Synthesis of Al-enriched γ’-phase coatings by pack cementation ........................... 141
5.3.2 Microstructure of the as-received coatings........................................................... 144
5.3.3 Isothermal oxidation performance of three Pt-diffused coatings ............................ 147
5.3.3.1 Elemental diffusion of three coatings ............................................................... 147
5.3.3.2 Oxide microstructure and growth kinetics......................................................... 150
LIST OF CONTENTS
4
5.3.3.3 TGO spallation .............................................................................................. 155
5.3.4 Rumpling behaviour of three bond coats under cyclic oxidation ............................ 156
5.4 Discussion ................................................................................................................ 161
5.4.1 Pt and Al depletion of three coatings ................................................................... 161
5.4.2 Effect of bond coat composition on the selective oxidation of aluminium................ 164
5.4.3 Rumpling behaviour of three coatings ................................................................. 165
5.4.3.1 Balint and Hutchinson rumpling model ............................................................. 165
5.4.3.2 B&H model applied to the Al-enriched γ’-phase coating...................................... 166
5.5 Summary ................................................................................................................. 167
Chapter 6 Conclusions and Future Work................................................................................ 169
6.1 Conclusions .............................................................................................................. 169
6.2 Future work ............................................................................................................. 171
Reference ......................................................................................................................... 173
Word counts: 43169
LIST OF FIGURES
5
List of Figures
Fig. 1.1 Progress in the maximum allowable temperatures of Ni-base superalloys and thermal
barrier coating (TBC) since 1965. The red line indicate the sharp increase of the allowable gas
temperature by the employment of TBCs [2]. .....................................................................19
Fig. 1.2 Schematic view of a TBC system on an airfoil [4]. .................................................20
Fig. 2.1 Schematic illustration of multi-layered and multifunctional TBC system. The functions
and properties for each layer are indicated [2]. ....................................................................26
Fig. 2.2 Coefficients of thermal expansion (CTEs) of a range of materials are cross-plotted
against their thermal conductivities [3]. ..............................................................................28
Fig. 2.3 Cross-sectional microstructure of a) APS and b) EBPVD TBC [24]. .......................30
Fig. 2.4 Binary phase diagram of Ni-Al system [39]. ..........................................................34
Fig. 2.5 The cross-sectional SEM images of the β-NiAl bond coats fabricated by a) low-activity
and high-temperature; b) high-activity and low-temperature pack cementation [41]. ............36
Fig. 2.6 The cross-sectional SEM image of as-fabricated (by pack cementation) β-NiPtAl bond
coat on the CMSX-4 superalloy substrate. ..........................................................................37
Fig. 2.7 The typical microstructure of a Pt-diffused γ-Ni/γ’-Ni3Al bond coat on the CMSX-4
superalloy substrate. The γ’-phase: brighter contrast; γ-phase: grey contrast. .......................39
Fig. 2.8 Cross-sectional images of the as-deposited NiCoCrAlY bond coat deposited by HVOF
[71]: a) optical image and b) back scattered electron (BSE) image (high magnification) showing
the β (grey contrast) + γ (white contrast) two-phase microstructure. The black contrast areas
are interfacial pores between metal particles. ......................................................................42
Fig. 2.9 A fractured cross-sectional image of a specimen exposed for several hours at 1200℃
showing the TGO columnar and equiaxed grains [4]. ..........................................................44
Fig. 2.10 a) Schematic illustration of the PLPS technique and b) typical R1/R2 fluorescence
spectra for Cr-containing stress-free (dashed line) and stressed α-Al2O3 (solid line) [81]. .....46
Fig. 2.11 The morphology of oxide scale on the alloy after h oxidation at ℃. Note the needle
or plate shape of θ-Al2O3 [88]. ...........................................................................................50
Fig. 2.12 Alloying elements in the Ni-based superalloys (adapted from [101]). ....................51
LIST OF FIGURES
6
Fig. 2.13 Microstructure of a Ni-based single crystal superalloy revealing a high volume
fraction of γ’ phase [100]: the cuboidal γ’ precipitates (grey contrast) in the γ-matrix (white
contrast). ...........................................................................................................................52
Fig. 2.14 Five major categories of failure mechanisms documented for TBC systems [74]. ..55
Fig. 2.15 The optical images of a EBPVD TBC sample showing a) the incipient buckling of
the top coat (viewed under reflected light) and b) subsequent spallation of the top coat [3]. ..57
Fig. 2.16 A schematic illustration of four primary cracking modes in an APS TBC system [11].
.........................................................................................................................................59
Fig. 2.17 SEM micrographs showing damage evolution in an APS TBC: a) isothermal
oxidation and b) thermal cycling [60]. ................................................................................59
Fig. 2.18 The cross-sectional SEM micrograph of an EBPVD TBC on a β-NiPtAl bond coat
exhibiting the TGO rumpling after 50 1-h cycles at 1150 ℃ [117]. .....................................61
Fig. 2.19 NiPtAl specimens showing the effect of the oxide thickness on the rumpling. The
oxide layer thickness is a) ~ 5 μm and b) ~ 10 μm. The systems were subjected to the same
thermal cycling history and obviously more rumpling developed for the specimen with thicker
oxide layer [127]. ..............................................................................................................63
Fig. 2.20 Schematic of different failure modes for a thin coating system. a) A single through-
thickness crack in the coating which deflects to the interface to cause the coating failure; b)
multiple through-thickness cracks in the coating; c) edge-delamination at the interface and d)
buckling-induced delamination for a compressed film [133]. ...............................................66
Fig. 3.1 a) Luminescence spectrum showing characteristic peaks for θ-Al2O3 and α-Al2O3 of
an oxide scale; b) R peaks of an α-Al2O3 scale on a Pt-diffused γ/γ’ coating after isothermal
oxidation at 1100°C for 1 h. For comparison, the spectrum from a stress-free, polycrystalline
alumina is shown by the red line. .......................................................................................76
Fig. 3.2 Cross-sectional SEM images of as-diffused samples: a) 2 µm Pt coating; b) 5 µm Pt
coating. .............................................................................................................................79
Fig. 3.3 Typical luminescence spectrum of a) 0 Pt, b) 2 µm Pt and c) 5 µm Pt samples after 2
min oxidation at 1000 ºC....................................................................................................80
Fig. 3.4 Luminescence spectrum of a) 0 Pt, b) 2 µm Pt and c) 5 µm Pt samples after 10 min
oxidation at 1000 ºC. .........................................................................................................81
LIST OF FIGURES
7
Fig. 3.5 Fraction profiles of θ-Al2O3 as a function of oxidation time. ...................................81
Fig. 3.6 a) - c): FIB/SEM cross-sectional images after 2 min oxidation; d) - f): FIB/SEM cross-
sectional images after 10 min oxidation at 1000 ºC of no Pt sample, 2 µm Pt coating and 5 µm
Pt coating, respectively. .....................................................................................................82
Fig. 3.7 a) - c): FIB/SEM cross-sectional images; d) - f): SEM surface images of no Pt sample,
2 µm Pt coating and 5 µm Pt coating after 30 min oxidation at 1000 ºC. ..............................84
Fig. 3.8 Oxide scale thickness evolution of no Pt sample and 5 µm Pt sample as a function of
oxidation time. ..................................................................................................................85
Fig. 3.9 The TGO stress evolution of no Pt sample and 5 µm Pt sample as a function of
oxidation time. ..................................................................................................................86
Fig. 3.10 a) ADF STEM image of 2 µm Pt sample after oxidation at 1050 ºC for 10 min; b)
combined phase map and phase reliability map obtained from automated crystal orientation
mapping in TEM, taken from the red box region in a). Green: θ-Al2O3; red colour: α-Al2O3.
.........................................................................................................................................87
Fig. 3.11 a) ADF STEM image of 5 µm Pt sample after oxidation at 1050 ºC for 10 min; b)
combined phase map and phase reliability taken from the red box region in a). Green colour:
θ-Al2O3; red: α-Al2O3. .......................................................................................................88
Fig. 3.12 Microstructure of the as-received a) Ni-20Al, b) Ni-20Al-10Pt, c) and d) Ni-20Al-
20Pt alloy. The inset in d) shows the magnified morphology of the tiny γ channels in the γ/γ’
region................................................................................................................................91
Fig. 3.13 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al alloy after oxidation at 1050 ºC for 2 min and 10 min, respectively. .......92
Fig. 3.14 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al-10Pt alloy after oxidation at 1050 ºC for 2 min and 10 min. ...................92
Fig. 3.15 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al-20Pt alloy after oxidation at 1050 ºC for 2 min and 10 min. ...................93
Fig. 3.16 a) Previously reported growth model of α-Al2O3: uniformly nucleate along the
interface, resulting in a layered structure of the oxide scale; b) and c) new growth model in our
study which illustrates lower Pt content and higher Pt content coatings, respectively. ...........95
LIST OF FIGURES
8
Fig. 4.1 Schematic view of the experimental setup: both high speed cameras take images of the
TBC coating surface (x-y plane) during the compression test, and the compression load is along
the y-direction. ................................................................................................................ 113
Fig. 4.2 As-received TBCs with different substrates, a) - d): cross-sectional SEM images; e)
and f): Ni, Al and Pt concentration profiles by SEM/EDX linescan along the lines in c) and d).
....................................................................................................................................... 115
Fig. 4.3 YSZ lifetime for the cyclic oxidation testing of TBC-coated different superalloy blades.
....................................................................................................................................... 117
Fig. 4.4 Cross-sectional SEM images: a) b) secondary electron (SE) mode; c) d) backscattered
electron (BSE) mode; e) f) corresponding SEM/EDX linescan elemental concentration profile
along the red lines in c) and d), respectively after 5 1-h cycles at 1200 °C.......................... 118
Fig. 4.5 a) b): cross-sectional SEM images (SE mode) and c) d): Ni, Al and Pt concentration
profiles by SEM/EDX linescan along the red lines in a) and b), respectively after 10 1-h cycles
at 1200 °C. ...................................................................................................................... 119
Fig. 4.6 a) and b): bond coat surface BSE images exposed by spalling after 28 cycles at 1200 °C;
c) and d): BSE images of the back side of spalled YSZ coating after 28 cycles. ................. 121
Fig. 4.7 a) TGO thickness evolution during the cycling test; b) root mean square roughness
evolution of TGO/BC interface by processing of digitized profiles. ................................... 122
Fig. 4.8 a) Optical image taken by the camera for DIC showing the sample surface prior to
applying the load. The region-of-interest (ROI) is highlighted with the red rectangular and is
used as the reference image f (x); b) deformed images at several loading scales corrected by a
DIC displacement field g (x + u (x)) and c) the corresponding residual field Φ (x) of a sample
(as-received TBC with N5 substrate) during the test. The red ellipse in c) highlights the
occurrence of buckling. .................................................................................................... 124
Fig. 4.9 Evolution of a) average buckling radii, and b) corresponding strains calculated by DIC
as a function of oxidation time at 1150 °C for specimens with the X4 and N5 substrate,
respectively. .................................................................................................................... 125
Fig. 4.10 a) Residual stress of TGO and b) average TGO thickness as a function of isothermal
oxidation time at 1150 °C for TBCs with the X4 and N5 substrate, respectively. ................ 126
Fig. 4.11 Mode Ⅰ interface toughness of TBCs on N5 (black square) and X4 substrates (red
circle) as a function of oxidation time, respectively. Data of other TBC systems are also
LIST OF FIGURES
9
included for comparison: Pt diffusion bond coat (1150 °C, X. Wang et al. [132]); NiCoCrAlY
bond coat (Yu-Fu Liu et al. [150]); β-NiPtAl bond coat (Vasinonta and Beuth [146])......... 129
Fig. 4.12 High resolution STEM/EDX analysis of the TGO/bond coat interface: a) and b)
HAADF (high angle angular dark field)/STEM image, b) is the red box area shown in a); c) -
g) STEM/EDX mapping of TGO on N5 substrate after 3 cycles at 1200 °C. ...................... 131
Fig. 4.13 High resolution STEM/EDX of the TGO/bond coat interface: a) and b)
HAADF/STEM image, b) is the red box area shown in a); c) - g) STEM/EDX mapping of TGO
on X4 substrate after 3 cycles at 1200 °C.......................................................................... 131
Fig. 5.1 a) cross-sectional FIB/SEM image and b) surface image of as-etched samples; c)
etching time-thickness plot and d) cross-sectional FIB/SEM image after Pt electroplating on
the etched substrate.......................................................................................................... 137
Fig. 5.2 Two steps to fabricate the Al-enriched γ’-phase coatings: Ⅰ. Fabricate Pt-diffused
intermediate coatings; Ⅱ. Pack cementation aluminizing on the intermediate coatings. ....... 139
Fig. 5.3 Ni-Pt-Al phase diagram at 1150°C [202]. The compositions of sample 1-5 are marked
by the different dots, respectively. The inset SEM image shows the as-fabricated cross-sectional
microstructure of sample 3. .............................................................................................. 143
Fig. 5.4 Ni-Pt-Al phase diagram at 1150°C [202]. The two horizontal red lines represent the
upper and lower limit of Al concentration for a pure γ’-phase coating, respectively. ........... 144
Fig. 5.5 XRD patterns of as-fabricated Al-enriched γ’-phase coatings, Pt-diffused γ/γ’ and β-
NiPtAl coatings. .............................................................................................................. 145
Fig. 5.6 Microstructure of as-fabricated coatings: SEM (BSE) images of: a) γ’ coating, b) γ/γ’
coating and c) β-NiPtAl coating; EBSD phase contrast map of the red box area in a - c: d) γ’
coating, e) γ/γ’ coating and f) β-NiPtAl coating; and corresponding color-coded inverse pole
figure (IPF) mapping g) - i) showing the different grain sizes of three coatings. ................. 147
Fig. 5.7 Ni, Pt and Al concentration evolution of Al-enriched pure γ’ coating by EDX linescan
after a) 0 h, b) 20 h and c) 50 h oxidation.......................................................................... 148
Fig. 5.8 Ni, Pt and Al concentration evolution of Pt-diffused γ/γ’ coating by EDX linescan: a)
0 h, b) 20 h and c) 50 h oxidation. .................................................................................... 149
Fig. 5.9 Ni, Pt and Al concentration evolution of β-NiPtAl coating by EDX linescan: a) 0 h, b)
20 h and c) 50 h oxidation. ............................................................................................... 150
LIST OF FIGURES
10
Fig. 5.10 Glancing angle (3°) XRD patterns of the oxides on a) γ’ coating, b) γ/γ’ coating and
c) β-NiPtAl coating after different oxidation time at 1150 °C. ........................................... 151
Fig. 5.11 a) - c) cross-sectional SEM images of three coatings after 20 h oxidation at 1150 °C;
d) the magnified image of the red box area in c)................................................................ 152
Fig. 5.12 Cross-sectional SEM backscattered electron (BSE) images of three coatings after 50
h oxidation at 1150 °C: a) - c) low magnification; d) - e) high magnification. ..................... 154
Fig. 5.13 Oxide thickness vs. isothermal oxidation time (at 1150 °C) for the three bond coat
systems. .......................................................................................................................... 154
Fig. 5.14 Optical surface images of three coatings after different isothermal oxidation time at
1150 °C. .......................................................................................................................... 156
Fig. 5.15 Profilometer images of (a) as-fabricated Al-enriched γ’-phase coating surface and
after (b) 5 10-minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the
identical area. (f) The surface profiles and the corresponding Rq of the line shown in (a). ... 158
Fig. 5.16 Profilometer images of (a) as-fabricated Pt-diffused γ/γ’ coating surface and after (b)
5 10-minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the identical area.
(f) The surface profiles of the line shown in (a). ................................................................ 159
Fig. 5.17 Profilometer images of (a) as-fabricated β-NiPtAl coating surface and after (b) 5 10-
minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the identical area. (f)
The surface profiles of the line shown in (a). .................................................................... 160
Fig. 5.18 The local interdiffusion flux of Al (𝐽𝐴𝑙) for the Al-enriched γ’-phase coating/CMSX-
4 diffusion couple and the Pt-diffused γ/γ’ coating/CMSX-4 diffusion couple at specific
positions (𝑥𝑖= 1, 5, 10 and 20 μm) after 20 h oxidation. The arrows represent the diffusion
direction. ←: uphill diffusion from inner part to the coating surface; →: from coating to the
inner part of superalloy. ................................................................................................... 164
LIST OF TABLES
11
List of Tables
Table 2-1 Comparison between APS and EB-PVD .............................................................29
Table 2-2 The structural properties of transient alumina phases ...........................................47
Table 2-3 Physical properties of Ni-based superalloys.........................................................53
Table 3-1 Composition of CMSX-4 substrates ....................................................................74
Table 3-2 Electroplating platinum bath ...............................................................................74
Table 3-3 Electroplating time and corresponding Pt thickness, average Pt concentration and
surface roughness ..............................................................................................................78
Table 3-4 θ-Al2O3 fraction of γ/γ’ areas and γ’ areas, respectively for the three Ni-Pt-Al alloys
after 2 min oxidation at 1050 ºC .........................................................................................93
Table 3-5 θ-Al2O3 fraction of γ/γ’ areas and γ’ areas, respectively for the three Ni-Pt-Al alloys
after 10 min oxidation at 1050 ºC .......................................................................................93
Table 3-6 Low-index (hkl) planes and the corresponding d-spacing values of two Ni phases (γ’
and γ) and θ-Al2O3: the d-spacing mismatch between planes of θ-Al2O3 and the corresponding
γ’ (or γ) plane with closest d-spacing matching is calculated and listed as the strain .............97
Table 4-1 Superalloy compositions (atomic %) by EDXRF ............................................... 110
Table 4-2 Elastic moduli of YSZ top coat measured by the cross-sectional nano-indentation
test .................................................................................................................................. 126
Table 5-1 Composition of CMSX-4 superalloy ................................................................. 137
Table 5-2 Electrolyte γ-etching bath parameters................................................................ 138
Table 5-3 Different pack cementation parameters and the resulting coating composition by
EDS after vacuum anneal................................................................................................. 142
LIST OF ABBREVIATIONS
12
List of Abbreviations
APS atmosphere plasma spray
BC bond coat
BSE backscattered electrons
CMAS calcium magnesium alumina silicate
CTE coefficient of thermal expansion
DIC digital image correlation
EDS energy dispersive X-ray spectroscopy
HAADF high angle annular dark field
PLPS photoluminescence piezospectroscopy
SE secondary electron
SEM scanning electron microscope
TBC thermal barrier coating
TEM transmission electron microscope
TGO thermally grown oxide
XRD X-ray diffraction
YSZ yttria-stabilised zirconia
LIST OF PUBLICATIONS
13
List of Publication
C. Liu, X. Zhang, Y. Chen, et al., Effect of superalloy substrate on the lifetime and interfacial
toughness of electron beam physical vapour deposited thermal barrier coatings, Surface &
Coatings Technology (2019), https://doi.org/10.1016/ j.surfcoat.2019.124937.
ABSTRACT
14
Abstract
Study of Bond Coats and Failure Mechanisms of Thermal Barrier Coatings
Chen Liu A thesis submitted to the University of Manchester for the degree of
Doctor of Philosophy, 2019
Bond coats for thermal barrier coating (TBC) applications and the failure mechanisms of TBCs
are addressed in this thesis, with a focus on i) the early stage oxidation of Pt diffusion bond
coats, ii) substrate effects on TBC failure and iii) a new bond coat design.
The early stage oxidation behaviours of γ/γ’-based NiAl bond coats with different Pt additions
are investigated. Pt can slow down the θ-Al2O3 to α-Al2O3 transformation. High resolution
phase mapping by scanning diffraction analysis shows that α-Al2O3 nucleation in the θ-Al2O3
scale is inhomogeneous along the coating/scale interface. Spatially resolved PLPS
(photoluminescence piezospectroscopy) results show a clear correlation between the θ-Al2O3
to α-Al2O3 transition and the γ or γ’ microstructure in the underlying alloy: where Pt stabilises
the γ’ structure, the suppression of θ-Al2O3 to α-Al2O3 transition is observed. The slower θ-
Al2O3 to α-Al2O3 transition rate due to Pt addition leads to a lower compressive stress of the
stable oxide scale, which contributes to the long term stability of the oxide scale.
The effects of substrate composition on the lifetime of TBCs were studied by comparing two
TBCs applied to a CMSX-4 and a René N5 single crystal superalloy substrate, respectively.
Both TBCs were applied by EB-PVD on top of the Pt-diffused γ/γ’ bond coats. Cyclic oxidation
test showed that TBCs deposited on the CMSX-4 substrates exhibited an average lifetime 20%
higher than that deposited on the René N5 substrate. The TGO thickness evolution and the
roughness of the TGO/bond coat interface were comparable for the two TBCs during cyclic
oxidation. To find out the mechanism for the substrate composition effect, a strain-to-fail test
combined with 3D-DIC was employed to measure the bond coat/TGO interface toughness and
its evolution for the two TBCs. The mode I interfacial toughness (Γic) values were almost
identical for the two TBCs (~ 30 J/m2) in the as-deposited state. However, it decreased much
faster for the TBC with a René N5 substrate after oxidation. The fast decrease of interface
toughness was attributed to the sulphur segregation at the TGO/bond coat interface.
A new Al-enriched γ’-Ni3Al bond coat has been developed and its high temperature oxidation
behaviour has been examined and compared with that of the conventional Pt-diffused γ/γ’
coating and the β-NiPtAl coating. This new γ’-phase coating exhibited significantly reduced
Al and Pt depletion during oxidation compared to the two conventional diffusion coatings.
Moreover, although the Al-enriched γ’-phase coating presented faster thermal grown oxide
(TGO) growth than that of the β-NiPtAl coating during isothermal oxidation, it outperformed
the β-NiPtAl coating regarding the rumpling resistance during cyclic oxidation. The Al-
enriched γ’-phase coating also exhibited superior TGO spallation resistance compared to the
Pt-diffused γ/γ’ coating. The mechanisms for the combination of good rumpling resistance and
oxidation performance of this γ’-Ni3Al coating will be addressed.
DECLARATION
15
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 other institute of learning.
COPYRIGHT STATEMENT
16
Copyright Statement
i) 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 s/he has given The University of
Manchester certain rights to use such Copyright, including for administrative purposes.
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ACKNOWLEDGEMENTS
17
Acknowledgements
I would like to express my gratitude to Professor Ping Xiao for offering me this chance to
conduct this research and acting as my mentor over the past years. His enthusiasm for the work
has inspired me and I will always be grateful for all the guidance, suggestions and
encouragement from him. I also thank Prof. Philip Withers, my co-supervisor, for all the
suggestions on my research.
I would like to thank my senior colleague, Dr. Ying Chen, who has guided me and helped me
since the very beginning of this project. Thanks to his expertise in the research of thermal
barrier coatings, I have learned a lot from him through all the discussions we had.
A special appreciation is given to Dr. Alexander S. Eggeman for conducting all the scanning
diffraction analysis (ASTAR) work shown here. This appreciation is extended to Dr. Xun
Zhang and Dr. Lin Qiu for the helpful suggestions on analysing the 3D-DIC data and
conducting the pack cementation experiments, respectively. I also would like to thank the
experimental officers in the School of Materials including Mr. Matthew Smith, Ms. Xiangli
Zhong, Dr. John Warren, Mr. Kenneth Gyves, Mr. Gary Harrison, Mr. Andy Wallwork, Mr.
Stuart Morse, Mr. Andrij Zadoroshnyj, Mr. Ben Spencer and Mr. Andy Forrest for the help on
my experiments.
I want to thank all the members in the ceramic coating group and all my friends in Manchester
for the help on my study and life.
I would like to thank my beloved families, my parents, Mr. Anwen Liu and Mrs. Yunchuan
Chen, and my grandparents for their love and support.
CHAPTER 1 INTRODUCTION
18
Chapter 1 Introduction
1.1 Gas-turbine engine materials
The gas-turbine engines used for the electricity generation and aircraft propulsion are Carnot
engines where their energy efficiency is proportional to the operating temperature of the turbine
[1]. However, the operating temperature of gas-turbine engines is restricted by the melting
point and high temperature properties of the turbine blade materials. The Ni-base superalloy
has been used as an exclusive structural material for the gas-turbine engines for the sake of its
high melting point (~ 1300 ℃) and superior mechanical properties at high temperatures. The
past decades have witnessed the development of the alloy composition design (for a
combination of oxidation resistance and creep resistance), directionally solidified and single
crystal Ni-base superalloys and internal cooling technologies [2, 3]. All of these have allowed
a steady increase in operating temperatures of gas-turbine engines over decades, which is
shown by the brown line in Fig. 1.1 [2].
A pursuit of higher engine efficiency in industrial applications has required an even higher
operating temperature of the engine. While a big hurdle to achieve this is that the current
operating temperature of the engines has exceed the melting point of the Ni-base superalloys
by ~ 200 - 300 ℃. This suggests that further improvement of the operating temperature is
unlikely to be achieved by optimising cooling technologies or alloy composition design alone.
Thermal barrier coatings (TBCs) were first proposed in 1980s to overcome this hurdle and
achieve higher operating temperatures of the engines. As shown by the red line in Fig. 1.1 [2],
the employment of TBCs has significantly increased the engine’s operating temperature, thus
increasing the engine efficiency.
CHAPTER 1 INTRODUCTION
19
Fig. 1.1 Progress in the maximum allowable temperatures of Ni-base superalloys and thermal
barrier coating (TBC) since 1965. The red line indicates the sharp increase of the allowable gas
temperature by the employment of TBCs [2].
1.2 Introduction of thermal barrier coatings
TBC is a complex, multi-layered and multifunctional system. A typical TBC system on an
airfoil is shown in Fig.1.2 [4], which consists of four layers: a ceramic top coat, a thermally
grown oxide (TGO) layer, a metallic bond coat and a superalloy substrate. The ceramic top
coat which is typically composed of 7-8 wt. % yttria stabilized zirconia (YSZ) is deposited on
the metallic bond coat. It can act as the thermal insulator due to its significantly low thermal
conductivity. The primary function of the metallic bond coat (which is firstly applied to the
substrate) is to provide a compatible bonding between the metallic substrate and the YSZ top
coat. More importantly, it also offers an oxidation resistance for the superalloy substrate by
forming a dense and protective TGO layer during high temperature oxidation. The wide
application of TBCs together with the state-of-art cooling technologies have effectively
CHAPTER 1 INTRODUCTION
20
lowered the surface temperature of the superalloy components, thus significantly increasing
the engine efficiency and prolonging the lifetimes of gas-turbine engines.
Fig. 1.2 Schematic view of a TBC system on an airfoil [4].
Although the use of TBCs have allowed improvements in the operating temperature, there are
still some problems and challenges with the current TBCs raised by further demands in better
durability and reliability of TBC systems [5]. Firstly, a prominent concern with TBCs is the
loss of adhesion and spallation from the underlying metal during service [6]. The underlying
bare metal would be directly exposed to the hot gas environment if the TBC spalls, which is
catastrophic. Hence, it is vital to build a reliable lifetime prediction model of TBCs based on a
comprehensive understanding of the failure mechanism of TBCs. Many phenomena including
phase transformation, oxidation, diffusion, sintering and thermal & mechanical deformation
occur concurrently in a TBC system during high temperature exposures. Therefore, it is of great
difficulty to describe the evolution of microstructures, stresses and thermal & mechanical
CHAPTER 1 INTRODUCTION
21
properties of TBCs during service and to further elucidate the failure mechanisms based on
these descriptions. For instance, despite years of research efforts, the effect of bond coat and
substrate chemistry on the degradation mechanism leading to early spallation of TBCs has not
been well understood [7]. To clarify this chemical composition effect, in-depth investigations
and analysis on the TBC system regarding various aspects (e.g. interfacial microstructure
evolution, oxide growth kinetics, etc.) are required.
Second, since bond coat is the key component in a TBC system and plays a crucial role in
controlling durability of TBCs, new bond coat design has attracted increasing attention to meet
the demand in higher temperature and higher efficiency. Future bond coat design should
consider three aspects for the sake of an improved performance of TBCs [8]:
i) The layer compatibility; the bond coat serves as an intermediate compliant layer to maintain
the top coat adhesion to the metal substrate.
ii) The bond coat must have adequate oxidation properties to form and maintain a protective
TGO scale (usually α-Al2O3).
iii) The mechanical properties; an ideal bond coat should exhibit superior high temperature
strength to resist deformation and is also ductile at room temperatures to be strain tolerant.
The combination of these three aspects makes ideal bond coat designs but it is not
straightforward due to the complexity and interactions of the TBC system.
Third, with increasing operating temperature in advanced engines, TBCs are increasingly
suspected to degrade by molten salts such as calcium-magnesium-alumina-silicate (CMAS) [9],
which have posed challenges for new TBC materials and engineering processing techniques.
CHAPTER 1 INTRODUCTION
22
1.3 Objectives and structure of the dissertation
The primary objective is to investigate the bond coat and substrate chemical composition effect
on the oxidation behaviour of TBCs in terms of TGO phase transformation, stress evolution,
surface rumpling and how these phenomena are correlated with the compositional changes of
the bond coat/substrate. In addition, a new bond coat design which combines the three
principles described in Section 1.2 has been proposed and successfully implemented. Detailed
descriptions for each chapter are listed as follows:
In Chapter 1, a brief introduction of gas-turbine materials and TBC system is given, followed
by the challenges in current TBC studies.
In Chapter 2, different components in a TBC system are reviewed in respect of material
selection, material properties and the processing methods. The degradation and failure
mechanisms for different TBC systems are also reviewed.
In Chapter 3, the effect of platinum addition on the early stage oxidation behaviour of Pt-
diffused γ-Ni/γ’-Ni3Al bond coats is investigated. Many previous studies have reported that Pt
can greatly improve the prolonged oxidation performance of the γ-Ni/γ’-Ni3Al bond coats.
However, only limited work has focused on the Pt effect on the early stage oxidation behaviour,
which can also affect the bond coat oxidation performance. Moreover, some controversial
views regarding Pt effect on the early stage oxidation of Pt-diffused bond coats have been
reported, so further study is required to clarify this issue. In this study, the early stage oxidation
behaviours of three Pt-diffused γ-Ni/γ’-Ni3Al bond coats with different Pt additions were
CHAPTER 1 INTRODUCTION
23
studied and compared in terms of the TGO phase transformation, residual stresses of TGO and
the TGO growth rate. Pt effect on the TGO phase transformation during early stage oxidation
is summarized and discussed. A new phase transformation mechanism based on high resolution
phase mappings by scanning diffraction analysis is proposed to explain the Pt effect on the
TGO phase transition.
In Chapter 4, the substrate composition effect on the lifetime of TBCs is studied by comparison
of two TBC systems with different superalloy substrates. Although a number of studies have
reported that the superalloy substrate composition can affect the cyclic lifetime of TBCs, the
mechanism of this substrate composition effect has not been fully understood. In this study,
two TBCs with Pt-diffused γ-Ni/γ’-Ni3Al bond coats are applied to a CMSX-4 and a René N5
single crystal superalloy substrate, respectively. Cyclic oxidation tests were carried out on these
two TBCs to compare their lifetimes as well as microstructural evolution and TGO growth
kinetics. A strain-to-fail test combined with 3D-DIC (digital image correlation) was employed
to measure the bond coat/TGO interface toughness and its evolution for the two TBC systems.
Finally, the mechanism of the substrate composition effect on the interface toughness evolution
and TBC lifetime was explored.
In Chapter 5, a new Al-enriched γ’-Ni3Al bond coat was designed according to the three
principles for the future bond coat design (as described in Section 1.2). This new bond coat
was deposited on the CMSX-4 superalloy substrate by selective etching of the CMSX-4
substrate combined with the low-temperature pack cementation. The isothermal oxidation
behaviours of this new Al-enriched γ’-Ni3Al bond coat have been studied and compared to the
CHAPTER 1 INTRODUCTION
24
conventional industry-standard Pt-diffused γ/γ’ coating and β-NiPtAl coating in terms of TGO
microstructure & growth rate, Pt & Al depletion and TGO spallation resistance. Furthermore,
thermal cycling to a maximum temperature of 1150 °C was conducted on all coatings to
compare their rumpling behaviours. The mechanisms for the good combination of rumpling
resistance and oxidation performance of this new γ’-Ni3Al-base coating are discussed.
In Chapter 6, the main conclusions and future work are summarised.
CHAPTER 2 LITERATURE REVIEW
25
Chapter 2 Literature Review
2.1 Thermal barrier coating system
The application of TBCs in conjunction with external cooling technologies can achieve a
temperature gradient up to ~ 170 ℃ across the coating [10]. As a result, the inlet gas
temperature and the efficiency of gas turbine engines have been significantly improved. A
state-of-the-art TBC system consists of four layers. Each layer and its functions and properties
are shown in Fig. 2.1 [2]. The nickel-base superalloy substrate is the structural component
which can sustain creep and cyclic fatigue. A bond coat layer is coated on the superalloy
substrate, which can provide adhesion between the superalloy substrate and ceramic top layer.
The TGO layer (usually α-Al2O3) forms between the metallic bond coat and the ceramic top
coat due to bond coat oxidation during high temperature exposures. The topmost layer is a
ceramic layer which acts as the actual heat shield. It usually has a combination of low thermal
conductivity and high strain tolerance. Due to the coupled diffusional and mechanical
interactions between each layer, the TBC system is dynamic during service and all layers
interact with each other to control the performance and durability of the system [11].
2.2 Top coat
2.2.1 Material requirement and selection
The first and foremost requirement for the top coat material is the low thermal conductivity
because its primary function is to provide the thermal insulation for the underlying metal
components [3]. In addition to the low thermal conductivity, the extreme thermomechanical
working environment has led to other requirements for the top coat material:
CHAPTER 2 LITERATURE REVIEW
26
(1) The top coat should have phase stability during prolonged thermal exposure because the
disruptive volume change accompanied with the phase transformation can cause damage of the
material [5].
(2) The top coat material must be stable to resist sintering and erosive pollutant attack (e.g.
CMAS) during high temperature exposures in an oxidising environment [12].
(3) The top coat should be strain-tolerant and fracture-resistant during thermal cycling. The
strain tolerance requires the material to withstand the strains induced by the thermal misfit
between the top coat and the metallic substrate during thermal cycling, and the fracture
resistance can mitigate the deformation during thermal cycling and impact damage from
airborne particles.
(4) The material should have thermodynamic compatibility with the TGO material (usually
alumina) to ensure a good interfacial adhesion of the TBC system.
Fig. 2.1 Schematic illustration of the multi-layered and multifunctional TBC system. The
functions and properties for each layer are indicated [2].
CHAPTER 2 LITERATURE REVIEW
27
All of these requirements have provided the guidelines for the top coat material selection and
development in the last decades and have laid the foundation for state-of-the-art TBC materials.
Currently, the top coat is typically composed of ~ 7 wt. % yttria stabilised zirconia (7YSZ)
because of its desirable properties which coincides with the above requirements for top coat
materials. First and foremost, as can be seen from Fig. 2.2 [3], YSZ has a low thermal
conductivity (~ 1-3 W/mK). This can be ascribed to its high concentration of point defects
including solute cations and oxygen vacancies, which can reduce the lattice thermal
conductivity by scattering lattice waves [13]. Moreover, despite being a ceramic material, the
coefficient of thermal expansion (CTE) of YSZ is well matched to that of the Ni-based
superalloys (Fig. 2.2), leading to reduced thermal misfit strains during thermal cycling
compared to other ceramics.
As described above, one of the requirements for the top coat material is the phase stability upon
thermal cycling to avoid volume changes associated with phase transformations that can lead
to the degradation. The dopants, such as Y2O3, MgO, CeO2, CaO and Sc2O3 can significant ly
enhance the phase transformation resistance compared to pure zirconia. Despite the fact that
all of these dopants can retard the phase transformation of ZrO2 to some extent, a large number
of studies have found that ZrO2 doped with 7 wt.% Y2O3 (7YSZ) exhibited the longest thermal
cycling lifetime and therefore is the most suitable candidate for the top coat [11]. Under
equilibrium condition, 7YSZ is stabilised as the tetragonal phase above 1050 ℃. This
tetragonal YSZ is transformed into a mixture of monoclinic and cubic YSZ upon cooling unless
mechanically constrained [14, 15]. However, due to the non-equilibrium fabrication process of
the top coats (e.g. electron beam physical vapour deposition), as-deposited 7YSZ coatings
typically have a metastable tetragonal prime (t’) phase instead of the equilibrium tetragonal (t)
phase [16]. Although these two phases are similar in structure, the t’ phase has been considered
CHAPTER 2 LITERATURE REVIEW
28
to be a non-transformable phase because it does not undergo any phase transformation after
prolonged thermal exposure at 1200 ℃ [17]. This makes the 7YSZ a good choice for TBC
applications.
As a refractory ceramic, YSZ has long-term sintering resistance during the high temperature
oxidation. In addition, the porous microstructure of TBC top coat resulting from the
manufacture process can contribute to the strain tolerance of the coating. Another merit of YSZ
is the high fracture toughness due to the ferroelastic toughening [18] combined with the phase
transformation toughening (martensitic transformation from tetragonal to monoclinic phase)
[19]. Thus the YSZ can fulfil the third requirement as listed above. Finally, YSZ exhibits
chemical compatibility and strong bonding with the alumina, which ensures the long-term
stability of the TBC system [3].
Fig. 2.2 Coefficients of thermal expansion (CTEs) of a range of materials are cross-plotted
against their thermal conductivities [3].
CHAPTER 2 LITERATURE REVIEW
29
2.2.2 Deposition techniques and microstructure
Currently, TBC top coats are generally deposited by plasma spraying (PS) or by electron beam
physical vapour deposition (EBPVD). PS uses a high-temperature plasma jet, to melt and
accelerate the powder feedstock (the material to be deposited) toward the substrate. The plasma
jet can be generated by passing a gas (usually Ar, He or N2) through an electric arc, during
which the gas will be ionised and form the plasma jet. Then the powder feedstock is injected
into the plasma jet either internally or through an external feed-port [20]. After the
instantaneously melting of the powder feedstock, the semi-molten powders are rapidly
accelerated towards the cold substrate. A mechanically bonded coating is deposited
immediately on the substrate surface by spreading and solidification. Depending on the process,
PS can be done at ambient conditions (atmospheric plasma spray, APS) or at controlled
conditions (e.g. vacuum). EBPVD was first introduced to fabricate TBCs in 1980s [21].
Electron beam is designed to transform target atoms into the gaseous form. These atoms then
precipitate into solid form and coat onto the preheated substrate [22]. Coatings deposited by
EBPVD are mainly used for extreme thermomechanical working conditions (e.g. blades of
aeroengines), while APS are more commonly used nowadays because of its operation
robustness and economic viability compared to EBPVD [23]. Table 2.1 gives a comparison of
these two coating techniques.
Table 2-1 Comparison between APS and EB-PVD
Methods Occasions [12] Interfacial bonding Equipment cost [12]
APS
Stationary parts on aeroengines;
stationary and rotating parts of
land-based power generation
engines
Mechanical bonding £0.6-1.1 million
EBPVD Vanes or blades of aircraft
engines Chemical bonding £10-20 million
CHAPTER 2 LITERATURE REVIEW
30
As mentioned above, due to the non-equilibrium fabrication process of the top coats (e.g.
EBPVD and APS), as-deposited 7YSZ coatings typically have a metastable tetragonal prime
(t’) phase instead of the equilibrium tetragonal (t) phase. In addition, APS and EBPVD TBCs
have totally different microstructures. As a result, they exhibit different advantages in terms of
properties. The thickness of APS TBCs is about hundreds of micros to several millimetres. Fig.
2.3 a shows that APS TBCs are featured by splat grains and inter-splat plate-like pores which
are parallel to the substrate/coating interface. These inter-splat pores resulting from the rapid
solidification can effectively reduce the thermal conductivity. On the other hand, because the
pores are parallel to the interfaces, APS TBCs generally have less strain compliance and
therefore shorter thermal cycling lifetimes than EB-PVD TBCs.
EBPVD TBCs are usually ~ 120-150 μm thick and exhibit columnar grain morphology with
inter-column gaps (Fig. 2.3 b). The gaps between the disconnected columns can accommodate
the thermal misfit strain between ceramic and metallic components during cycling exposures
and thus providing better strain compliance. However, the lack of large splat pores normal to
the heat flow direction can lead to higher thermal conductivity compared to APS TBCs [12].
Fig. 2.3 Cross-sectional microstructure of a) APS and b) EBPVD TBC [24].
CHAPTER 2 LITERATURE REVIEW
31
2.2.3 New top coat candidates
The upper limit of use temperature for 7YSZ is 1200 ℃. This can be explained by two reasons.
First, the t’ phase of YSZ will decompose into a high yttria cubic phase and a low yttria
tetragonal phase after the long term exposure at elevated temperatures [14]. The latter will
transform to a monoclinic phase upon cooling accompanying with a large volume expansion,
which accelerates the TBC failure. Another reason for the temperature limit is the sintering of
YSZ at temperatures above 1200 °C leading to loss of strain tolerance and early failure [23].
Therefore, considerable efforts are being invested in identifying new top coat materials with
better performance than the current industry-standard 7YSZ [23, 25-27].
Several studies found that some rare earth element doped zirconia, such as pyrochlores
(A2B2O7) and perovskites (ABO3) have lower thermal conductivity than 7YSZ [23].
Furthermore, their thermal stability and sintering resistance are better compared to 7YSZ,
which makes them promising candidates for the TBC top coat. However, relatively lower CTE
(e.g. YSZ CTE: 11×10-6 °C, La2Zr2O7 CTE: 9×10-6 °C [27]) of these materials impedes their
development because the thermal mismatch between the substrate and coating can induce
higher thermal stresses. Recent studies [26, 28] have proposed a double ceramic top coat
combining YSZ with another ceramic of lower thermal conductivity, which exhibited longer
thermal cycling life than the single YSZ layer. Currently, since no single material can meet all
the requirements for the extremely complicated TBC system, the double layer top coat seems
promising in future TBC applications.
CHAPTER 2 LITERATURE REVIEW
32
2.3 Bond coat
2.3.1 Material requirements
The YSZ top coat is not deposited directly onto the substrate. Since YSZ is transparent to
oxygen diffusion, during service, oxygen would diffuse through the interconnected pores of
YSZ to the interface and oxidize the superalloy by forming fast growing Ni-rich oxides [29].
These Ni-oxides (e.g. NiO and Ni(Cr, Al, Co, Ti, Ta)2O4) are thermodynamically incompatible
to the YSZ layer and can cause early spallation of the top coat [30]. To avoid this problem,
bond coats with sufficient high temperature oxidation resistance are applied to the superalloy
substrates prior to the TBC deposition.
The bond coat functions as an intermediate adhesion layer between the YSZ top coat and the
superalloy substrate. In addition to this, the most significant function of bond coats is to provide
sufficient oxidation resistance for the superalloy substrates by forming a slow-growing and
protective TGO scale on the surface of the bond coat when it oxidizes at high temperatures.
Currently, the industry-standard bond coats are commonly made of alloys with specific
aluminium-rich compositions, which can result in the formation of a TGO scale that mainly
consists of α-Al2O3. The dense α-Al2O3 scale can prevent the oxidation of the underlying
superalloy and is also thermodynamically compatible with the YSZ top coat. Apart from the
oxidation resistance, the bond coat is also expected to minimize the interdiffusion with the Ni-
based superalloy substrate [31]. Furthermore, the bond coat material should be morphologically
stable to resist plastic deformation of the surface induced by thermal stresses during thermal
cycling. Because the surface deformation can cause interfacial debonding and spallation failure
of the coating.
CHAPTER 2 LITERATURE REVIEW
33
2.3.2 Bond coat categories
There are two categories of bond coats based on the fabrication techniques: diffusion coatings
and overlay coatings [32]. The former can be further classified into two categories by their
phase constituents. One is comprised primarily of the β-NiAl phase (referred to as nickel
aluminide coatings), and the other is composed of γ-Ni/γ’-Ni3Al phase. The overlay coatings
are typically made of MCrAlY (M=Ni, Co or a combination of both) alloys. The following
sections will elaborate the deposition technique, composition, microstructure and properties of
these coatings, respectively.
2.3.2.1 β-NiAl based diffusion bond coat
The β-NiAl based diffusion bond coats are fabricated by a high temperature interdiffusion
annealing process between an external aluminium source and the nickel superalloy substrate.
The external aluminium source can be applied by vapour deposition methods such as vapour-
phase aluminizing (VPA) [33], chemical vapour deposition (CVD) [34, 35] and pack
cementation [36, 37]. These methods all involve in a reaction between an aluminium donor and
a halide activator firstly to generate the gaseous aluminium halide. Then this gaseous
aluminium halide reacts with the nickel superalloy substrate and forms the β-NiAl based nickel
aluminide coatings at elevated temperatures. Among these deposition methods, the pack
cementation has been widely applied due to its low cost, processing simplicity and flexibility
for different specimen dimensions and geometries [38]. For the pack cementation process, the
pack powder, which consists of the aluminium source (usually pure Al), an inert filler (e.g. α-
Al2O3) and a halide activator such as NH4Cl or AlCl3, is ground and mechanically mixed firstly.
Then the specimen is buried in the pack powders charged into an air-tight alumina crucible.
The crucible is then heat treated in a protective atmosphere (e.g. argon gas). During the heat
treatment, two procedures take place simultaneously in the pack [38]. One is the
activation/migration process, which includes the chemical reaction to create the aluminium
CHAPTER 2 LITERATURE REVIEW
34
halide (AlXn, X=F, Cl or Br and 1≤n≤3) vapour and the migration of the vapour to the specimen
surface. Another process (termed as deposition/diffusion process) is the deposition of
aluminium by the oxidation-reduction reaction between the aluminium halide and the metal,
followed by the interdiffusion between the deposited aluminium and alloying elements (e.g.
Ni) in the metal substrate. The key step is the generation of the aluminium halide vapour. This
gaseous phase will react with the metal to deposit aluminium by the following reaction:
𝐴𝑙𝑋𝑛 +𝑛
𝑛 − 1𝑁𝑖 → 𝐴𝑙 +
𝑛
𝑛 − 1𝑁𝑖𝑋𝑛−1
The β-NiAl phase has a B2 crystal structure, which is an ordered body centred cubic (bcc)
structure consisting of two simple cubic interpenetrating sublattices. From the Ni-Al binary
phase diagram (Fig. 2.4 [39]), it can be seen that the β-NiAl phase composition range is wide,
indicating that the Al concentration of β-phase can vary significantly from its NiAl
stoichiometric composition.
Fig. 2.4 Binary phase diagram of the Ni-Al system [39].
CHAPTER 2 LITERATURE REVIEW
35
The microstructure of the β-NiAl coating is relied on the pack cementation process.
Specifically, there are two categories of pack cementation according to the heat treatment
temperature: the low-activity and high-temperature (above ~ 1000 ℃) pack cementation and
the high-activity and low-temperature (below ~ 950 ℃) pack cementation [40]. The first type
(low-activity) coating has two zones: the outer zone (Zone 1 in Fig. 2.5 a) and inner zone (Zone
2 in Fig. 2.5 a). Both zones have a nickel-rich β-NiAl-phase matrix. Zone 1 contains various
amounts of substrate alloying elements (e.g. Cr, Mo, Co, and Ti) in solution, and Zone 2 has a
variety of dispersed phases including MC and M23C6 carbides and σ(Cr, Mo, Co) phases [41].
While the high-activity coating has three zones: the outer Zone 1, the middle Zone 2 and the
inner Zone 3 (Fig. 2.5 b), and all of them have an Al-rich β-NiAl-phase matrix. Specifically,
Zone 1 (Fig. 2.5 b) contains dispersed α(Cr, Mo) phases and carbides, and the latter are identical
to those found in the underlying Ni-based superalloy [41]. The middle Zone 2 (Fig. 2.5 b) is
comprised of σ-phase and carbides in the β-NiAl-phase matrix, which is similar to the outer
Zone 1 of the low-activity coating. And Zone 3 (Fig. 2.5 b) also shows similar microstructure
to that of Zone 2 for the low-activity coating (Fig. 2.5 a).
CHAPTER 2 LITERATURE REVIEW
36
Fig. 2.5 The cross-sectional SEM images of the β-NiAl bond coats fabricated by a) low-activity
and high-temperature; b) high-activity and low-temperature pack cementation [41].
The distinct microstructures of these two types of β-NiAl coatings can be explained by the
coating formation mechanisms. The formation of low-activity coatings are based on
predominant outward diffusion of Ni. On the other hand, the high-activity coating forms as a
result of the predominant inward diffusion of Al, and a higher amount of Al can diffuse into
the coating during the aluminizing process due to the significantly higher Al activity. This can
explain the Al-rich β-phase matrix for the high-activity coating. Another characteristic
difference is that the low-activity coating has much less amount of complex precipitates
throughout the coating, especially in the outer zone. However, the high-activity coating has
numerous precipitates in the entire coating. This is due to the fast inward diffusion of Al during
coating formation. The slow diffusing alloying elements (e. g. Ta, Mo etc.), which are
originally from the substrate, may become trapped in the fast-forming β-NiAl matrix, and
precipitate out due to their limited solubility in the β-phase [42].
CHAPTER 2 LITERATURE REVIEW
37
The β-NiAl phase, due to its high Al concentration, can serve as an Al reservoir for the
formation and continuous growth of the protective Al2O3 scale during thermal exposures. This
dense alumina scale can protect the underlying superalloy from being oxidized. However,
during long-term oxidation, especially under cyclic conditions, the alumina scale can spall off,
which results in the loss of YSZ top coat and failure of the TBC system. Thus, the TGO scale
adhesion is significant for evaluating the oxidation performance of the bond coat. Pt addition
is the most widely applied method to improve the adhesion of alumina scales [42-44]. The
state-of-art industrial process to fabricate the platinum modified nickel aluminide (β-NiPtAl)
coatings includes a Pt electroplating process (5-7 μm Pt layer), followed by an aluminizing
process by pack cementation or CVD [5]. The microstructure of the as-fabricated β-NiPtAl
bond coat is similar to that described for the β-NiAl coating, as shown in Fig. 2.6.
Fig. 2.6 The cross-sectional SEM image of the as-fabricated β-NiPtAl bond coat on the CMSX-
4 superalloy substrate.
CHAPTER 2 LITERATURE REVIEW
38
The β-NiPtAl coatings exhibit significantly improved TGO scale adhesion and spallation
resistance, compared to the unmodified β-NiAl coatings. For instance, Hou and Tolpygo [44]
have conducted cyclic oxidation tests at 1150°C in air for both NiAl and NiPtAl coatings on
the same single crystal superalloy substrate. The TGO scale on the β-NiAl coating spalled after
~ 300 10-min cycles, while the oxide scale of the β-NiPtAl coating remained adherent even
after 2000 10-min oxidation cycles. Numerous studies [42, 45-47] have reported that Pt
addition can improve the TGO adherence, and some possible mechanisms have been proposed
to explain this Pt effect. One possibility is that Pt addition can directly strengthen the interface
bonding between the alumina and the NiAl alloy. However, Svensson et al. [48] have calculated
the work of separation of the α-Al2O3(0001)/β-NiAl (111) interface, in pure and Pt-rich NiAl
materials respectively using the density functional theory (DFT). They found that the interfacial
bonding is decreased with Pt addition, which excludes the interfacial strengthening effect of Pt.
Some studies [44, 46] confirmed that platinum can reduce the interfacial pores. However, it is
uncertain whether this effect alone can fully explain the improved scale adhesion. It has also
been proposed that Pt can prevent the sulphur segregation at the TGO scale/coating interface
[45, 49], or limit the outward diffusion of minor alloying elements (e.g. Ta, Ti, Re etc.) from
the substrate [45], thus improving scale adhesion. In a word, currently there is no well-accepted
mechanism for this Pt effect on the TGO scale adhesion.
2.3.2.2 Pt-diffused γ-Ni/γ’-Ni3Al bond coat
The success of Pt addition to the β-NiAl bond coats has inspired the invention of another
diffusion coating, the Pt-diffused γ-Ni/γ’-Ni3Al bond coat. Since a problem associated with the
β-NiPtAl bond coats is that the conventional aluminizing process (e.g. pack cementation) can
introduce undesirable elements or impurities that can lead to poor adhesion between the bond
coat and the top coat. The Pt-diffused γ-Ni/γ’-Ni3Al bond coats have eliminated the
CHAPTER 2 LITERATURE REVIEW
39
aluminizing process by electroplating a thin layer of Pt (7 - 12 μm) on the single crystal
superalloy substrate followed by a diffusion heat treatment at 1100 -1200 ℃ up to several hours
[50]. This not only reduces the fabrication cost, but also increases the coating stability because
the γ-Ni/γ’-Ni3Al phases have good chemical compatibility with the superalloy substrate [51].
The Pt-diffused γ-Ni/γ’-Ni3Al bond coats are formed by the interdiffusion between the Pt layer
and the nickel-based superalloy. During the process, the inward diffusion of Pt into the
superalloy will destroy the original γ/γ’ lattices of the superalloy and expand the unit cells due
to the large atom radius of Pt. The new γ’ and γ phases precipitate out with Pt in solid solution
and exhibit a strip-like microstructure. After extended heat treatment, the strip-like
microstructure is elongated, as shown in Fig. 2.7. The Pt-enriched γ-Ni phase has a fcc (face
centred cubic) structure and the Pt-enriched γ’-Ni3Al phase has a L12 structure. The γ’-phase
shows a brighter contrast with more Pt enriched in the Z contrast image.
Fig. 2.7 The typical microstructure of a Pt-diffused γ-Ni/γ’-Ni3Al bond coat on the CMSX-4
superalloy substrate. The γ’-phase: brighter contrast; γ-phase: grey contrast.
CHAPTER 2 LITERATURE REVIEW
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Platinum plays a crucial role in the oxidation performance of the γ/γ’-based bond coat. First,
Pt in the γ/γ’-based bond coat has sustained the benefit of improving oxide scale adhesion as
described in Section 2.3.2.1 for the β-NiAl bond coat. Second, Pt can promote the selective
oxidation of aluminium. In other words, Pt can promote the growth of the protective Al2O3
scale and inhibit the growth of detrimental Ni-oxides such as NiO or spinel. According to Izumi
and Gleeson [52], the extents of NiO and spinel formation decreased significantly with
increasing Pt content during cyclic oxidation of γ/γ’-based nickel alloys at 1150 ℃, which in
turn improved the oxidation resistance of the γ/γ’ alloys. Briefly, there are three factors
contributing to the selective oxidation of aluminium due to Pt addition [53, 54]:
(1) The inert nature of Pt can ensure that no Pt involved oxidation reaction takes place even at
extremely high temperatures (~ 1500 ℃).
(2) The subsurface enrichment of Pt at the TGO/metal interface can reduce the chemical
activity of aluminium (aAl) at the interface, causing an uphill diffusion of aluminium from the
inner part of the metal to the surface, which promotes the exclusive growth of the Al2O3 scale.
(3) Pt replaces almost solely to the Ni sites in the γ/γ’ structure, which increases the Al:Ni atom
ratio on a given crystallographic plane containing both Ni and Al. As a result, the aluminium
oxidation is favoured rather than nickel.
Apart from the Pt contribution, a recent study [55] has pointed out that numerous grain
boundaries near the surface of the Pt-diffused γ-Ni/γ’-Ni3Al bond coat can provide fast
diffusion paths of aluminium at initial stage of oxidation, which also contributes to the
exclusive growth of Al2O3 scale.
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As mentioned above, the Pt-diffused γ-Ni/γ’-Ni3Al bond coat has some advantages over the β-
NiPtAl bond coat including low fabrication cost and better phase stability. Furthermore, this
bond coat also exhibits negligible morphological instability during cyclic oxidation due to its
higher creep resistance compared to the β-based coating [56]. However, a concern with this
bond coat is its relatively low aluminium concentration, leading to the formation of detrimental
Ni-oxides after long-term high temperature oxidation. These Ni-oxides such as spinel can
significantly degrade the oxide/coating interface adhesion due to their brittleness, which leads
to premature failure of the coating system [57]. A new bond coat design by improving the
aluminium concentration while maintaining the structure of γ’-phase seems to be attractive for
optimizing the TBC system. Actually, some efforts have been made to improve the aluminium
concentration of Pt-diffused γ-Ni/γ’-Ni3Al bond coats by a secondary aluminizing process
(pack cementation) [58, 59]. However, the as-fabricated Al-enriched γ-Ni/γ’-Ni3Al bond coats
only exhibited limited oxidation performance improvement due to the Al depletion during
oxidation. Therefore, further studies on this issue are demanding for optimizing the current
TBC system.
2.3.2.3 MCrAlY overlay bond coat
The MCrAlY overlay bond coats are directly sprayed onto the surface of Ni-based superalloy
substrates using physical deposition techniques such as APS [60, 61], high velocity oxygen
fuel spraying (HVOF) [55, 62, 63], low-pressure plasma spraying (LPPS) [12, 64] and EBPVD
[65]. Unlike the diffusion coatings, the overlay coating composition and thickness can be
tailored by the coating source and deposition time with great flexibility and accuracy, which
makes this coating a desirable candidate in some applications.
The compositions of MCrAlY (M=Ni, Co or a combination of both) overlay bond coats are
typically: (in wt. %) ~ 15 - 25 % Cr, ~ 10 - 15 % Al, ~ 0.2 - 1% Y and Ni (Co) in balance [12].
CHAPTER 2 LITERATURE REVIEW
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Typical cross-sectional images of a NiCoCrAlY bond coat deposited by HVOF are shown in
Fig. 2.8. From Fig. 2.8 a), it can be seen that this deposition technique can produce a dense
coating with the uniform thickness and low oxide content. The as-deposited NiCoCrAlY bond
coat exhibits a β (grey contrast) + γ (white contrast) two-phase microstructure (Fig. 2.8 b). The
β-phase can serve as the main aluminium source to improve its oxidation resistance. The TGO
scale on this bond coat is mainly composed of α-Al2O3. However, after prolonged high
temperature oxidation, some Y/Al rich oxides are also observed in the TGO scale apart from
α-Al2O3 [66]. On the other hand, the γ-phase is designated to improve its mechanical ductility.
Chromium is added to enhance the hot corrosion resistance of the coating [67]. In addition,
some studies also reported that Cr can promote the selective growth of α-Al2O3 by a third-
element effect mechanism [68], thus also contributing to the oxidation performance. The minor
additions of reactive element (RE) yttrium in this bond coat can significantly improve the TGO
scale adhesion thus extending the lifetime of the coating. Although several mechanisms have
been proposed to explain this RE effect, it is commonly accepted that yttrium can improve the
scale adhesion by segregating to the metal/scale interface thereby preventing the detrimental
sulphur segregation [69, 70].
Fig. 2.8 Cross-sectional images of the as-deposited NiCoCrAlY bond coat deposited by HVOF
[71]: a) optical image and b) back scattered electron (BSE) image (high magnification) showing
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the β (grey contrast) + γ (white contrast) two-phase microstructure. The black contrast areas
are interfacial pores between metal particles.
2.4 Thermally grown oxide
2.4.1 Material requirements and selection
The oxygen from the engine environment can diffuse through the YSZ top coat during the
service of gas turbine engines. As a result, the bond coat will oxidise and form a thermally
grown oxide (TGO; 1 - 10 μm in thickness) layer between the bond coat and the YSZ top coat
[72]. Two factors have contributed to the oxygen diffusion from the engine environment to the
bond coat surface during oxidation. First, the interconnected pores in the YSZ top coat have
provided diffusion paths for oxygen to the bond coat surface. Second, the high oxygen
diffusivity of the YSZ top coats makes it ‘oxygen transparent’ [11, 73].
The extreme operating conditions have raised several requirements for the TGO material:
1) The TGO material should be phase compatible with YSZ to ensure thermodynamic stability
of the top coat for long-term high temperature exposure [3].
2) It should be a slow-growing and stable oxide at high temperatures [3]. The fast-growing
oxides can result in very thick TGO layer during high temperature exposures, which increases
the trend for oxide spallation due to an increasing strain energy in the TGO scale.
3) It should have low oxygen diffusivity [74]. The TGO layer must perform as an oxygen
diffusion barrier to protect the underlying metallic part from being oxidised.
4) It should be mechanically robust to resist fracture, especially in highly cyclic scenarios.
Considering all of these requirements, the major classes of bond coats have been developed to
form a TGO scale which is predominantly composed of α-Al2O3 during high temperature
exposures in air.
CHAPTER 2 LITERATURE REVIEW
44
2.4.2 TGO microstructure and stress
The pure Al2O3 TGO layer usually exhibits two microstructural zones after prolonged oxidation:
a columnar zone (CZ, columnar grains) next to the bond coat and an equiaxed zone (EZ,
equiaxed grains) next to the YSZ (Fig. 2.9 [4]). The inner columnar oxide grains are formed
by the inward diffusion of oxygen anions and the outer equiaxed grains are formed by outward
diffusion of aluminium cations.
Fig. 2.9 A fractured cross-sectional image of a TGO scale showing the columnar grains formed
by the inward diffusion of oxygen and equiaxed grains formed by the outward diffusion of
aluminium [4].
Generally, the residual stress of TGO at room temperature consists of two components: the
thermal mismatch stress and the growth stress. The former is generated in the TGO scale by
cooling from the elevated temperature to the ambient temperature due to the thermal expansion
mismatch between the TGO and the Ni-based superalloy [75]. Since α-Al2O3 has a much lower
coefficient of thermal expansion (CTE; 8 - 9 × 10-6 m·C-1 [74]) than that of the Ni-based
superalloy (12 - 16 × 10-6 m·C-1), the thermal mismatch stress is usually compressive (~ -3 - 5
GPa). The growth stress is generated mainly due to the lateral growth of TGO scale [76]. The
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45
lateral elongation of TGO is constrained by the underlying metal substrate thus generating a
compressive growth stress in the scale.
While the thermal mismatch stresses can be calculated quite well, the magnitude of the growth
stress is determined by the dynamic competition between two opposite processes at high
temperatures: the stress generation and concurrent creep relaxation of the TGO layer.
Consequently, the magnitude of the growth stress is evolving during the oxidation process and
is dependent on the lateral growth strain, creep strain of the TGO layer and plastic deformation
of the underlying metal [77]. Several studies have reported the magnitude of growth stress of
alumina scales on various alloys. For instance, Schumann et al. [78] have measured the growth
stress of α-Al2O3 scales on NiAl alloys during oxidation at 1100 ℃ by X-ray diffraction, and
concluded that the compressive growth stress was very low (less than ~ 50 MPa). On the other
hand, the growth stress of alumina scales on FeCrAl alloys measured by the same technique
was ~ -1 GPa [79].
Knowledge of the residual stress of TGO is crucial because the strain energy associated with
the TGO stress and thickness results in the delamination of the TGO scale from the underlying
metal [4]. In practice, the premature damage of the TGO scale is usually accompanied with
variations of local TGO stress (e.g. decrease in magnitude or change of stress state) in the
damaged area. Since the TGO is buried underneath a thick YSZ top coat, there is a need for
non-destructive tools to probe through the entire YSZ top coat and detect the TGO stress before
the failure occurs [80]. One promising tool is photoluminescence piezospectroscopy (PLPS).
PLPS technique can measure the TGO stress underneath the YSZ layer based on the R
luminescence (Fig. 2.10 [81]), which is generated by the phonon emission of chromium (Cr3+)
impurity in the Al2O3 scale when appropriately excited by an argon laser [82, 83]. Two
CHAPTER 2 LITERATURE REVIEW
46
fluorescence transitions of the dopant ion Cr3+ in the Al2O3 crystal correspond to the R1 and
R2 fluorescence doublet as shown in Fig. 2.10. Since the fluorescence transition is extremely
sensitive to the local ionic environment in the host crystal, stresses (or deformations) which
alter the interionic distances will shift the position of the doublet. The linear relationship
between the frequency shift of R luminescence and the in-plane equi-biaxial TGO stress
magnitude has been calibrated with a precision of ~ 10 MPa by J. He and co-workers [84]. In
addition, other luminescence parameters are also good indicators for the damage quantification
in TBCs, such as the peak shape, intensity ratio of R1 and R2, peak broadening and peak
separation [80, 85, 86].
Fig. 2.10 a) Schematic illustration of the PLPS technique and b) typical R1/R2 fluorescence
spectra for Cr-containing stress-free (dashed line) and stressed α-Al2O3 (solid line) [81].
2.4.3 TGO transformation during early stage of oxidation
As mentioned above, α-Al2O3 is the predominant oxide in the TGO during the service of TBCs.
However, before the stable growth of the α-Al2O3 scale, one or more alumina polymorphs
CHAPTER 2 LITERATURE REVIEW
47
usually grow firstly during early stage of oxidation. These phases, also called transient
aluminas, will transform to the stable α-Al2O3 at higher temperatures or after longer exposures.
The following phase transformation sequence has been reported in the literature during the
early stage oxidation of alumina-forming alloys [87-90]:
γ 750℃→ 𝛿
900℃→ 𝜃
1000℃→ 𝛼
The structural properties of these transient alumina phases are summarized in Table 2.2. The
lattice parameters, space groups and the orientation relationships with respect to γ-Al2O3 of
these three transient alumina phases are listed. δ-Al2O3 has been described as a superlattice of
the spinel structure with ordered cation vacancies and has a tetragonal symmetry. θ-Al2O3 is
the most widely studied polymorph of alumina and has a monoclinic symmetry and space group
c2/m. The aluminium cations are equally distributed over the octahedral and tetrahedral sites.
Table 2-2 The structural properties of transient alumina phases
Phase Lattice parameters
[88]
Space group Cations/unit cell
[89,90]
Orientation
relationship
with respect to
γ-Al2O3 [89,90]
γ aγ= 7.9 Å Fd3̅m 64/3 -
δ aδ= 7.9 Å
bδ = 15.8 Å
cδ = 11.9 Å
P212121 64 (100)δ║ (100)γ
[100]δ║ [001]γ
θ aθ= 11.9 Å
bθ = 2.8 Å
cθ = 5.7 Å
β=104°
C2/m 8 (100)θ║ (001)γ
[010]θ║[110]γ
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48
Overall, these transient alumina phases are all based on a defective spinel structure. The
transient alumina polymorphs can be described as a fcc array of oxygen anions with aluminium
cations partially filling the tetrahedral and octahedral interstices [91]. For these alumina phases
with the defective spinel structure, the ratio of octahedral to tetrahedral sites occupied by Al3+
decreases from γ-Al2O3 (the ratio is 2) until θ-Al2O3 (the ratio is 1). In other words, the
transformations within the transient regime can be described in terms of a change in site
occupancy of cations [91]. For example, some researchers [91] have applied electron
diffraction to study the alumina transitions and concluded that both γ-Al2O3 and δ-Al2O3 were
based on the fcc packing of the oxygen anions but with a higher degree of order for the
interstitial cations in the δ phase. The stable α-Al2O3 is trigonal symmetry with rhombohedral
centring (space group R3c). The oxygen anions are hexagonal close packing (hcp) and the
aluminium cations occupy the octahedral sites in the anion sublattice [92].
The investigations considering the early stage oxidation of alumina-forming alloys are
intensively focused on the θ-Al2O3 formation, its transformation to α-Al2O3 and the effect of
additives on this phase transformation. For example, Prasanna et al. [88] have studied the effect
of θ-Al2O3 formation on the growth kinetics of FeCrAlY alloys during oxidation. They
demonstrated that θ-Al2O3 formation can significantly enhance the oxidation rate of alumina-
forming alloys. Furthermore, they have utilised a two-stage oxidation test which was performed
in 𝑂16 2 / 𝑂18 2 gas at 900 ℃ and secondary ion mass spectropy (SIMS) analysis to elucidate the
θ-Al2O3 to α-Al2O3 transformation process. The needle-shaped θ-Al2O3 (Fig. 2.11 [88]), which
grew by outward diffusion of aluminium, transformed to equiaxed small grains of α-Al2O3
which grew by inward diffusion of oxygen after early stage oxidation. Yang et al. [92] have
investigated the transient oxidation stage of single crystal (001) NiAl alloys by the electron
diffraction analysis and found that α-Al2O3 nucleated at the oxide/alloy interface with random
CHAPTER 2 LITERATURE REVIEW
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orientations. Other studies have investigated the transient oxide formation on NiAl alloys by a
variety of techniques including scanning electron microscopy (SEM) [53, 89, 91], transmission
electron microscopy (TEM) [90, 93], PLPS [94] and thermogravimetric analysis (TGA) [89,
95, 96].
The effect of different additives on the θ-Al2O3 to α-Al2O3 phase transformation is also of
interest because different additives can retard or accelerate this transformation thus affecting
the oxide growth kinetics. For instance, hafnium (Hf) can slow down the θ-Al2O3 to α-Al2O3
transformation and result in a consistent fast growth of oxide scale during the transient
oxidation stage of Pt-modified γ’-Ni3Al-based alloys [96]. Conversely, Brumm and Grabke [89]
concluded that Cr addition in the Ni-Al alloys can accelerate the θ-Al2O3 to α-Al2O3
transformation. This Cr effect is explained by the Cr2O3 nuclei formation in the initial stage of
oxidation which serves as nucleation sites for α-Al2O3, causing a faster transition rate. The
yttrium (Y) effect is more complex. According to Jedlinski [97], Y can accelerate or retard the
transient alumina to stable α-Al2O3 transformation depending on the amount and form of
yttrium in the alloy as well as its form in the oxide scale. Specifically, small amount of Y
accelerates the phase transition by provision of heterogeneous nucleation sites. While higher
amount of Y can be incorporated into the lattice of transient alumina by forming mixed Y-Al
oxides such as Y3Al5O12, and retard their transition into stable α-Al2O3. As for the Pt addition,
controversies exist regarding its effect on this phase transformation [94, 98, 99]. Unlike Cr or
Y, Pt is inert and cannot form any oxide or dope into the oxide scale in ion form. This suggests
that other mechanisms should be responsible for the observed Pt effects on this transformation.
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Fig. 2.11 The morphology of oxide scale on the Fe-20Cr-Al alloy after 72 h oxidation at 900 ℃.
Note the needle or plate shape of θ-Al2O3 [88].
In summary, different additions of alumina-forming elements can affect the transient alumina
to stable α-Al2O3 transition, which in turn has an influence on the oxide growth kinetics.
However, reliable mechanisms to explain the effect of additives are still demanding mostly due
to lack of microscopic information of the transient oxide distribution and α-Al2O3 nucleation.
2.5 Superalloy substrate
The Ni-based superalloys have been applied as the structural materials of aircraft and power-
generation turbine blades over the last few decades. After years of alloy development, the state-
of-the-art Ni-based superalloys can tolerate average temperatures ~ 1050 ℃ with occasional
local hot spots to temperatures ~ 1200 ℃, which is ~ 90% of the melting point of the Ni alloy
CHAPTER 2 LITERATURE REVIEW
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[100]. The following sections will briefly review the composition, microstructure and
characteristic properties of the Ni-based superalloys.
2.5.1 Composition and microstructure
The nickel-aluminium binary system is the basis for the Ni-based superalloy composition for
gas-turbine applications. Apart from this, at least five to ten alloying elements are also added
to the superalloy which can account for up to 40 wt. %. Typical alloying elements of the Ni-
based superalloys are given in Fig. 2.12 [101].
The fcc γ-phase is the basic constitute of the Ni-based superalloy. With increasing aluminium
contents, a precipitate γ’-Ni3Al phase forms, which has an ordered intermetallic L12 structure.
These two phases are the major constitutes of the Ni-based superalloy. Fig. 2.13 [100] shows
the microstructure of a Ni-based single crystal superalloy with a high volume fraction of
cuboidal γ’ precipitates in the grid-shape γ matrix. Single crystal Ni-based superalloys
including CMSX-4, SRR 99 and René N5 all show this typical microstructure with different
minor element additions, which will be illustrated below.
Fig. 2.12 Alloying elements in the Ni-based superalloys (adapted from [101]).
CHAPTER 2 LITERATURE REVIEW
52
Fig. 2.13 Microstructure of a Ni-based single crystal superalloy revealing a high volume
fraction of γ’ phase [100]: the cuboidal γ’ precipitates (grey contrast) in the γ-matrix (white
contrast).
Some alloying elements, such as W, Mo and Re are solid solution strengtheners of the γ-phase.
While other alloying elements including Ti, Ta and Nb can strengthen the γ’-phase by forming
the γ’- Ni3(Al, Ti, Ta, Nb) intermetallic compound [102]. For example, both the CMSX-4 and
SRR 99 single crystal superalloy contain Ti as strengtheners, while the René N5 superalloy
excludes this strengthener. Furthermore, Cr can improve the corrosion resistant of the
superalloy and Y can contribute to the oxidation resistance of the superalloy [103]. The trace
elements, e.g. C, B, Zr and Hf can form carbides or borides which are often located at the grain
boundaries. These elements are added to control the grain structure thus affecting the
mechanical properties that are strongly influenced by the grain boundaries [100].
2.5.2 Physical and mechanical properties
A large fraction of the structural components in gas-turbine engines are made of the Ni-based
superalloys for the sake of their exceptional combination of physical and mechanical properties.
Table 2.2 lists typical physical properties of Ni-based superalloys. In practical applications, it
CHAPTER 2 LITERATURE REVIEW
53
is usually worth considering the density normalised properties, especially for the rotating
components in the gas-turbine engines.
Table 2-3 Physical properties of Ni-based superalloys
Properties Typical ranges
Density 7.7 - 9.1 g/cm3 [100]
Melting point 1320 - 1450 ℃
Thermal conductivity 9 - 11 W/(m·K) (RT)
CTE 12 - 18 ×10-6 /℃ [104]
Ni-based superalloys exhibit relatively high yield tensile strength (~ 900 - 1300 MPa at room
temperatures). However, the tensile properties show a significant decay at temperatures above
850 ℃. The single crystal CMSX-3 superalloy exhibits a yield strength ~ 200 MPa at
temperatures of 1000 ℃ [105]. Modern Ni-based superalloys are optimised to improve the high
temperature creep resistance, which is crucial because during service the superalloys are under
stress for extended periods at high temperatures. Creep properties are influenced by the
alloying elements and microstructure of the superalloys. For example, the additions of
refractory elements Re, W and Mo have been successful in improving the creep resistance of
Ni-based superalloys. However, further incorporation of these refractory elements is limited
due to the formation of topologically close packed (TCP) phases at elevated temperatures,
which is associated with the initiation of creep damage. New generation superalloys have
incorporated Ru to suppress TCP formation resulting in enhanced creep properties [106].
2.6 The degradation and failure of TBCs
The core value of applying the YSZ top coat is to protect the underlying superalloy substrates
from the high temperature environment. Any degradation (erosion) or failure (spalling away
from the superalloy) of the top coat will lead to direct exposure of the superalloy to the hot
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gases with temperatures over its melting point, which is catastrophic. Therefore, establishing
reliable lifetime models and appropriately evaluating the TBC lifetime are very important for
the development of prime-reliant TBCs. The prerequisite for establishing reliable lifetime
models is to gain a comprehensive understanding of the failure mechanisms for various TBC
systems under different exposure conditions.
In general, TBC systems exhibit various failure mechanisms, and some well-accepted ones are
listed in Fig. 2.14 [74]. Based on the nature of these mechanisms, they can be summarised into
intrinsic and extrinsic mechanisms. The former is associated with thermal processes such as
elemental interdiffusion (i and ii in Fig. 2.14), TGO growth (iii in Fig. 2.14) and bond coat
deformation (iv). The extrinsic mechanism refers to the environmental degradation including
foreign object damage (FOD, v) and CMAS (calcium-magnesium-alumino-silicate) attack.
Specifically, in some cases, the alternative oxides such as spinels form either between the TGO
(α-Al2O3) and the bond coat due to Al depletion (i) or between the top coat and the TGO
because of the Ni diffusion through the TGO (ii) [10]. The brittle spinel can result in the
interfacial delamination. Alternatively, some failures are dominated by the strain energy
density of the oxide scale, and its interplay with the imperfections in the vicinity (non-planar
interface, iv) or within the TGO (planar interface with imperfections, iii) [6]. This failure
mechanism will be emphasized in this study and will be elaborated in Section. 2.6.1. In other
cases, the particle impact (FOD) may locally compress the porous top coat, leading to local hot
spots which can accelerate oxidation and contribute to the failure process [107]. Since the
extrinsic failure mechanism will not be involved in the next chapters, the following discussions
will only focus on the intrinsic mechanisms which will be organised as follows: firstly, a brief
introduction of the general failure modes of TBCs, followed by the failure mechanisms for
specific TBC systems under different exposure conditions.
CHAPTER 2 LITERATURE REVIEW
55
Fig. 2.14 Five major categories of failure mechanisms documented for TBC systems [74].
2.6.1 General failure modes
As mentioned above, the TGO develops a large compressive stress when cooling from the
service temperature. This stress in the thin TGO layer can cause the undulation (morphological
instability) of the scale, which can result in the failure of the system. This failure process can
be envisioned as a sequence of the crack nucleation, propagation and coalescence events [74,
108, 109]. First of all, tensile stresses which are normal to the bond coat/TGO interface are
CHAPTER 2 LITERATURE REVIEW
56
induced as a result of the TGO imperfections. This process is associated with an strain energy
release rate, thus initiating small cracks in the vicinity [6]. Then the stresses around the
imperfections and the associated energy release rates govern the propagation of these small
cracks or separations. The TBC remains attached at the remnant ligaments during the crack
propagation process. Finally, when cracks from the nearby TGO imperfections coalesce and
the remnant ligaments are detached, the TBC spalls either by edge delamination or large-scale
buckling (LSB) [74].
The above-mentioned two competing failure mechanisms: edge and buckle driven
delamination, both have been documented in previous studies [6, 110-112]. In this work the
emphasis will be on the buckling driven delamination, which is prevalent for TBC failures. Fig.
2.15 [3] shows a typical buckling driven failure of an EBPVD TBC system, which is common
for the compressive thin films. Consider a thin film subject to an equi-biaxial compressive
stress state, σ0, which is given by [110]:
𝜎0 = 𝐸∆𝛼∆𝑇/(1− 𝜈) (2.1)
where ∆𝑇 is the temperature drop from which the film is stress-free. E is the young’s modulus
of the film and ν is Poisson’s ratio. ∆𝛼 is the difference between the CTE of the substrate
(having a higher CTE, thus ∆𝛼 > 0) and the thin film. A buckling index is then defined as
[110]:
𝛱 = (1 − 𝜈2)(𝜎0/𝐸)(𝐿/ℎ)2 (2.2)
where h is the film thickness and L is the size of the separation which is present at the
film/substrate interface. When Π exceeds a critical value Πc (Πc =4.89 for a circular buckle
[110]), the buckling can be initiated. Equating Π=Πc =4.89 into Eq. (2.2), a critical separation
size Lb at the onset of buckling can be expressed as [110]:
𝐿𝑏/ℎ = 2.21√𝐸/𝜎0 (2.3)
CHAPTER 2 LITERATURE REVIEW
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where 𝐸 is the plane strain modulus of the film, 𝐸 = 𝐸/(1 − 𝜈2). Here, 𝐿𝑏 represents the
smallest separation size to initiate a buckle in absence of any TGO imperfection. By
substituting typical 𝐸 and 𝜎0 values into Eq. (2.3), Lb is ~ 20 ℎ in magnitude, which is about
several mm when considering a film thickness of ~ 100 μm. This relatively large size of flaw
can hardly be seen in the as-deposited TBCs. Hence, the initiation and growth of small
separations must occur until these separations reach the critical size needed for LSB. In addition,
the calculation of energy release rates of separations that are induced by the imperfections of
thin films have validated that TGO imperfections (e.g. undulations or heterogeneities) can
decrease the critical flaw size for the buckle initiation [110].
Fig. 2.15 The optical images of a EBPVD TBC sample showing a) the incipient buckling of
the top coat (viewed under reflected light) and b) subsequent spallation of the top coat [3].
In brief, the general failure mode for a TBC system includes a sequence of crack initiation,
growth & coalescence, and buckle driven delamination events. The TGO heterogeneities and
undulations play a key role in the failure process by re-distributing the stress in the vicinity of
the TGO imperfections, thus facilitating the above sequence of events. APS TBCs and EBPVD
TBCs are totally different in their microstructure and thermophysical properties, so different
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failure processes apply respectively [113]. The following sections will discuss the failure
process for the two TBC systems in detail.
2.6.2 The failure mechanisms of APS TBCs
As described in Section 2.2.2, YSZ top coats deposited by APS have inter-splat pores and
cracks which are parallel to the substrate/coating interface. In addition, since the bond coat
surface is roughened (the roughness average Ra can be tens of micro) before the deposition of
the YSZ layer to enhance the mechanical bonding, the metal/ceramic coating interface is highly
undulated. Because of the complex microstructure, the mechanics of events which proceeds
the edge- or buckle-driven delamination is still not fully understood [113]. The undulations at
the bond coat/TGO/top coat interfaces can give rise to out-of-plane tensile stresses which are
normal to the metal/ceramic interface. These stresses, combined with the interface
imperfections lead to the failure of the APS TBC system. Fig. 2.16 [11] gives a schematic
illustration of four primary cracking modes documented for APS TBCs. Thermal mismatch
stresses are developed at the TGO/bond coat interface upon cooling, and the interface
undulations can redistribute the thermal misfit stresses: the tensile stresses at the undulation
crests and the compressive stresses at the troughs [114]. The out-of-plane tensile stresses
increase as the TGO thickens, causing the crack initiation at the bond coat/TGO interface (type
Ⅰ crack in Fig. 2.16) of the crest. The type Ⅱ crack is initiated at the top coat/TGO interface of
the crest, and type Ⅲ corresponds to the cracking within the top coat in the vicinity of the crest.
These types of cracks are also resulted from the out-of-plane tensile stresses in the vicinity of
the TGO/top coat interface. The mechanism for type IV crack in Fig. 2.16 is more complicated.
When the TGO scale is very thin, the thermal mismatch stress between the bond coat (BC) and
the top coat (tbc) (determined by the CTE difference, αBC - αtbc >0) results in out-of-plane
tension in the top coat above the undulation crests and compression in the troughs. As the TGO
thickens, it constitutes a good fraction of the bond coat asperity, therefore the thermal mismatch
CHAPTER 2 LITERATURE REVIEW
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stress is locally dominated by the CTE mismatch between the bond coat/TGO ‘composite’ and
the top coat, rather than just the bond coat [115]. Beyond a certain critical TGO thickness, the
CTE of the bond coat/TGO ‘composite’ becomes lower than that of the top coat (αtbc), thereby
reversing the stress in the YSZ undulation troughs from compression to tension [11]. This
tensile stress drives the cracking within the troughs of the YSZ layer as shown by type IV.
Fig. 2.16 A schematic illustration of four primary cracking modes in an APS TBC system [11].
Lastly, the thermal treatment can affect the cracking mode for the APS TBCs [116]. According
to Trunova et al. [60], the isothermal heat treatment tends to promote crack propagation within
the TGO, while cyclic exposure can shift the crack path towards the top coat. Fig. 2.17 [60]
shows examples of different crack paths caused by the isothermal and cyclic oxidation.
Fig. 2.17 SEM micrographs shows the damage evolution in an APS TBC: a) isothermal
oxidation and b) thermal cycling [60].
CHAPTER 2 LITERATURE REVIEW
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2.6.3 The failure mechanisms of EB-PVD TBCs
Unlike the APS TBCs, the crack initiation of EBPVD TBC systems is mainly at the TGO/top
coat interfaces or the TGO/bond coat interfaces, rather than inside the TGO. Here, the failure
mechanisms of EB-PVD TBCs will be addressed in terms of different bond coat types.
The top coat deposited by EBPVD usually applies to the MCrAlY or β-NiPtAl bond coat on a
relatively flat surface (Ra is about a few micros). However, during prolonged thermal exposures,
especially thermal cycling, the bond coat surface is gradually roughened, due to the progressive
displacement of TGO into the bond coat, as shown in Fig. 2.18 [117]. This undulation that
develops in the TGO layer, also termed as rumpling, ratcheting or undulation instability, has
attracted considerable interest as a specific form of TBC failure with the MCrAlY or β-NiPtAl
bond coats [118-121]. Because the rumpling growth can generate tensile stresses across the top
coat/TGO/bond coat interfaces. These tensile stresses can initiate interfacial cracking between
the top coat and the superalloy substrate, leading to LSB and spallation failure of the TBC
[121].
Previous studies have extensively investigated the rumpling of TGO on the MCrAlY or β-
NiPtAl bond coats by experiments or simulations in the last decades [118, 121-124]. Among
these, Tolpygo and Clarke [121-123] have carried out systematic studies to find out the
mechanisms behind this rumpling behaviour. By conducting a series of variable-controlling
experiments on the surface roughness evolution of the nickel-aluminide bond coats, they have
ruled out some previously suggested factors that contribute to the rumpling and revealed that
the rumpling is mainly driven by a combination of TGO lateral growth and the thermal
mismatch between the coating and the underlying superalloy. For instance, they have
conducted thermal cycling above the martensitic transformation temperature Ms of the β-
CHAPTER 2 LITERATURE REVIEW
61
NiPtAl bond coat and observed the same roughness evolution compared to the thermal cycles
with martensitic transformation, which excluded the influence of martensitic transformation on
the rumpling [122]. They also pointed out that volume changes in the bond coat (due to Al
depletion induced phase transformation, e. g. β-NiAl to γ’-Ni3Al) are not necessary for the
rumpling initiation. Because the rumpling initiation can take place at very beginning of
oxidation before the phase transformation (due to Al depletion) occurs [121].
Fig. 2.18 The cross-sectional SEM micrograph of an EBPVD TBC on a β-NiPtAl bond coat
exhibiting the TGO rumpling after 50 1-h cycles at 1150 ℃ [117].
A mechanistic model (Balint & Hutchinson, B&H model [125]) has been developed to simulate
the rumpling behaviour of the TGO layer under various exposure conditions. This model can
provide an analytical approximation of the rumpling growth by considering factors including
TGO thickening and high temperature yielding of the TGO. The fundamental idea of this model
is that the rumpling deformation is driven by the lateral growth strain of TGO and assisted by
concurrent creep of the bond coat in compliance with the TGO deformation. The predictions
of rumpling growth on systems with the top coat and without the top coat have been presented
using this model [125]. For the TBC system with the top coat, more than a ten-fold reduction
of the total rumpling growth is predicted compared to that without the top coat. This prediction
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62
agrees well with the observations in the literature. For example, it is observed that the rumpling
is exclusively in the form of downward displacement of the TGO layer into the bond coat,
rather than both upward and downward displacements of TGO when no top coat is present
[126]. Because the top coat can significantly restrain the upward displacements of the TGO
and reduce the total rumpling growth. Furthermore, Balint and Hutchinson [125] have applied
this model to explore several important effects on the rumpling growth such as TGO thickness
effect and thermal history effect. The B&H model predicts more rumpling for the thicker oxide
layers, which is consistent with the experimental results by Tolpygo and Clarke (Fig. 2.19
[127]).
In addition to the TGO rumpling induced by thermal cycling, the MCrAlY bond coat has
another form of TGO imperfection which can also initiate cracking - the TGO thickness
heterogeneity due to the growth of yttrium-rich oxides (such as Y2O3 or yttrium aluminium
garnet, YAG) [128]. The TGO heterogeneities, also termed as ‘pegs’, can be developed and
enlarged in regions where the yttrium-rich oxides form which have a much higher oxygen
diffusivity than that of α-Al2O3. When this TGO heterogeneity develops to a critical size,
tensile stresses are generated around it followed by interfacial separations [74]. However, the
influence of this TGO heterogeneity on TBC failure is in debate. Other studies suggested that
these pegs can mechanically anchor the TGO layer in the bond coat, thus increasing the
interfacial adhesion and contributing to TBC lifetimes [128, 129].
CHAPTER 2 LITERATURE REVIEW
63
Fig. 2.19 NiPtAl specimens show the effect of the oxide thickness on the rumpling. The oxide
layer thickness is a) ~ 5 μm and b) ~ 10 μm. The systems were subjected to the same thermal
cycling history and obviously more rumpling developed for the specimen with thicker oxide
layer [127].
Unlike the MCrAlY or β-NiPtAl bond coat, the EBPVD TBCs based on the Pt-diffused γ/γ’
bond coats are not suspected to develop significant rumpling during cyclic exposure [86]. In
addition, the TGO on this bond coat did not exhibit thickness heterogeneity due to the
formation of fast-growing oxides [130]. It has been widely observed that for this TBC system,
the failure occurs mainly at the TGO/bond coat interface. Zhao and Xiao [131] have proposed
that during oxidation, impurities such as sulphur and refractory elements segregate to the bond
coat/TGO interface and degrade the TGO adhesion, which causes failure of the TBC. Other
researchers have found that the undulations of the initial non-planar bond coat surface can lead
to local interface separations due to the normal tensile stress across the interface [130]. These
separations gradually coalescence, causing spallation failure of the system. In summary, the
failure mechanism of EBPVD TBCs with the Pt-diffused γ/γ’ bond coat is still not fully
understood and remains to be further investigated.
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64
2.6.4 Interface toughness measurement of TBCs
The above sections have summarised different factors that can lead to crack initiation for the
APS TBC and the EBPVD TBC systems, respectively. In general, the lifetime of TBCs is
determined by the interaction between the driving force for delamination (the energy release
by crack initiation) and the resistance to crack propagation (toughness) [132]. Therefore, in
addition to the driving force of failure, the resistance force, i.e. the interface toughness is also
critical for understanding the failure mechanisms and establishing reliable lifetime models of
TBCs. This section will firstly introduce the definition of interface adhesion & delamination
and interfacial toughness, followed by some general failure modes for a thin film system.
Finally, some widely-used approaches to investigate the interface toughness for a thin ceramic
film on the metal substrate will be reviewed, which might be applied to the TBC systems.
2.6.4.1 Definitions of interface adhesion & delamination and interface toughness
The interface adhesion can be defined as: ‘the state in which two surfaces are held together by
interfacial forces including electrostatic forces, Van der Waals forces or chemical bonding’
[133]. Delamination, on the other hand, is the phenomenon that a coating separates from the
underlying substrate, which is driven by external or internal stresses, corrosion, etc.
The true work of adhesion for the interface is the amount of energy required to create free
surfaces from the bonded layers [134]:
𝑊𝐴 = 𝛾𝑓 + 𝛾𝑠+ 𝛾𝑓𝑠 (2.4)
where 𝛾𝑓 and 𝛾𝑠 are the specific surface free energies of the film and the substrate, respectively.
𝛾𝑓𝑠 is the specific free energy of the interface. The true work of adhesion is an intrinsic property
of the interface and can be determined by contact-angle experiments [135]. However, in
practice, the delamination is usually associated with energy dissipation due to different
CHAPTER 2 LITERATURE REVIEW
65
mechanisms (e. g. contamination, plastic deformation, etc.) during most of the adhesion tests.
Thus, the true work of adhesion is difficult to extract from the measurements. Instead, a
practical work of adhesion or the so-called critical strain energy release rate is defined as [133]:
𝐺𝑖𝑛𝑡 = − 𝜕𝑈
𝜕𝐴 (2.5)
where U is stored elastic energy released and A is the area of the interfacial crack. Then the
interface toughness, Kint can be defined by 𝐾𝑖𝑛𝑡 = √𝐸𝑖𝑛𝑡𝐺𝑖𝑛𝑡, where 𝐸𝑖𝑛𝑡 is a representative
Young’s modulus for the coating/substrate system.
The above definitions of the interface delamination and the critical strain energy release rate
provide the basic theories for interface toughness measurements. Then some general modes of
failure are introduced prior to the introduction of interface toughness tests. Fig. 2.20 outlines
four common failure types of a coating system [133]:
(a) A single through-thickness crack in the coating, and it propagates to the interface and
induces the coating delamination [136];
(b) A periodic array of through-thickness cracks which divert to the interface and induce the
failure [137, 138];
(c) Crack is generated at the interface and propagates along the interface [110];
(d) For a compressed coating, an initial crack at the interface can grow and lead to the buckling-
driven failure of the coating [132].
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66
Fig. 2.20 Schematic of different failure modes for a thin coating system. a) A single through-
thickness crack in the coating which deflects to the interface to cause the coating failure; b)
multiple through-thickness cracks in the coating; c) edge-delamination at the interface and d)
buckling-induced delamination for a compressed film [133].
Consider a well-developed interfacial defect of radius c, from a simple fracture mechanics
approach, the interface toughness 𝐾𝑖𝑛𝑡 is given by:
𝐾𝑖𝑛𝑡 = 𝑎𝜎√𝜋𝑐 (2.6)
where 𝑎 is the geometry constant which is ~ 1 for a circular defect. 𝜎 is the stress to cause the
coating delamination.
The above-mentioned energy dissipation during the delamination process depends on the mode
mixty angle (phase angle 𝛹), a measure of the relative amount of shear and normal stress
components at the crack tip [134], given by:
𝛹 = tan−1(𝐾Ⅱ/𝐾Ⅰ) (2.7)
where 𝐾Ⅱ and 𝐾Ⅰ are the interface toughness of mode Ⅱ failure (pure shear, Ψ = 90°) and mode
Ⅰ failure (pure opening fracture Ψ = 0°), respectively. The amount of energy dissipation
increases as the phase angle increases from 0° to 90°. The interface fracture toughness as a
CHAPTER 2 LITERATURE REVIEW
67
function of the mode mixty angle 𝛹 has been proposed by several studies, and the widely-used
one is defined by Hutchinson and Suo [139]:
𝛤(𝛹) = 𝛤0{1+ 𝑡𝑎𝑛2[(1 − 𝜆) ψ]} (2.8)
where 𝛤0 is the mode Ⅰ interface toughness for Ψ = 0°. λ is a parameter ranging from 0 to 1,
which is dependent on the interfacial friction.
2.6.4.2 Interface toughness test methods
In this section, several common tests to determine the interface toughness will be briefly
reviewed, including the indentation tests, pull-off tests, pushout methods, micro-cantilever tests
and four-point bending tests. The focuses for each method will be the experimental setup, the
strengths & limitations and its applications to the TBC systems.
Indentation tests. Indentation is the most common technique to measure the mechanical
properties of hard coatings such as Young’s modulus and hardness. It can also be used to model
the fracture behaviour, thus measuring the interface fracture toughness. Especially for thin
films, indentation is an effective method for sampling a small area of the interface with
common laboratory equipment for toughness measurement [140, 141]. Different indenters can
be selected according to the coating systems and the testing conditions, such as spherical
indenter (blunt), Vickers indenter (sharp), etc. The fracture modes are determined by the
indenter geometry and material properties. There are five common fracture modes for brittle
films including the cone cracks, Palmqvist radial cracks, median cracks, lateral cracks and the
half-penny radial cracks [142]. The corresponding fracture toughness analysis for each fracture
mode has been well established by the stress-based approach [143, 144] or the energy-based
approach [138, 142, 145]. Although the indentation test is featured by its experimental
simplicity and no requirement for the sample geometry, only limited work has been done to
test the toughness of TBCs by this method [132, 143, 146]. This can be attributed to the facts
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68
that indentation is highly dependent on the plastic deformation behaviour of the substrate and
the pores/cracks in the YSZ layer, both of which are hard to determine.
Barb pushout methods. The barb pushout method has been successful in evaluating the
interface toughness of EBPVD and APS TBCs [147]. The experimental setup is built on a ‘barb’
geometry, which originates from the pushout test for the fibre reinforced ceramic/metal matrix
composites [148]. The TBC top coat layer is notched at a distance of several mm from the end
of the specimen, and the remaining segment of the top coat is carefully removed by polishing.
Then two specimen pieces like this are affixed back-to-back, to form a ‘T-bone’ shape
specimen for the barb pushout test. This specimen is supported by two blocks during the
pushout test and the force-displacement curve is recorded. Then the steady-state energy release
rate can be simply estimated by the elastic properties of the YSZ top coat and the substrate,
combined with the maximum load and specimen dimensions. This methodology was initially
developed to analyse the crack developing process, where the crack propagation occurs
predominantly under the pure shear (mode Ⅱ) loading [149]. Liu et al. [150] have further
determined the phase angle of this loading condition to be ~ 65 - 70°. The benefit of this test is
that it is not sensitive to the anisotropic structure of YSZ top coats. While the main
disadvantage of this technique is the relatively complex sample preparation process.
Micro-cantilever tests. This method has been regarded as a particularly useful tool to measure
the mechanical properties and the interface toughness of coating systems with anisotropic
materials [151]. The micro-cantilever can be made by the focused ion beam (FIB) which is
usually coupled with SEM. Then the nanoindentation can be applied on the micro-cantilever
to deflect it until it snaps. The interface toughness can be expressed in terms of the maximum
applied load, the size of the notched crack and the dimensions of the cantilever beam based on
CHAPTER 2 LITERATURE REVIEW
69
the simple beam theory [152]. This method has been successfully applied to measure the
interface toughness of various ceramic coatings on the metallic substrates including the
SiO2/copper system [152], TaN/copper system [153] and the APS TBC system [154]. The main
disadvantage of this technique is the tedious sample preparation process because it requires a
symmetric cross-section of the cantilever beam with accurate dimensions to meet the simple
beam theory conditions.
Pull-off methods. There are two types of pull-off tests, (i) the tape peel test [141] and (ii) the
tangential lap shear test [155]. The tape peel test uses a piece of pressure-sensitive tape to peel
off the film. By measuring the force at which the film is removed, the peeling energy can be
obtained. This is a simple method but it can only be applied to the weak bonding coatings such
as polymers. In the tangential lap shear test, the coating is pulled off by an adhesively mounted
rod with a tensile force. This test can measure harder coatings. However, a perfect alignment
to ensure uniform loading across the coating/substrate interface is challenging for some
systems, especially those with special specimen shapes [133]. Thus, so far this method has
limited applications for evaluating the TBC systems.
Four-point bending. In this technique, a notch must be machined through the coating and
symmetric interfacial precracks are generated. Then the notched bending beam specimen is
monitored during the onset of the delamination with increasing load, and the interface
toughness can be simply calculated provided that there is no plastic deformation. For some
coating systems, the plastic deformation of the substrate can occur. So a stiffening layer is
attached to the coating system in order to avoid the plastic deformation, which is so-called
modified four-point bending test [156, 157]. The four-point bending test has been widely
utilised to test the interface toughness of ceramic coatings on softer substrates. However, the
CHAPTER 2 LITERATURE REVIEW
70
application of this technique is limited due to the following reasons. First, it requires special
specimen geometry and the pre-cracking can be difficult for some systems which involves
complex specimen preparation [141]. Second, in practice, the discontinuous crack development
of a non-planar crack front can occur, which diverts the failure from the computable mode.
Thus, only limited work has been done by this technique for the TBC systems. For instance,
Zhao et al. [157] have reported the energy release rate to be ~ 50 J/m-2 for the as-sprayed APS
TBCs by the modified four-point bending test.
To sum up, these tests all involve inducing coating delamination in a controlled process to
allow for a quantitative analysis for the interface toughness. The measured toughness values
can be affected by test-specific factors and the residual stress of the film [158]. The interface
toughness is a critical parameter for the TBC applications, which requires the analysis of failure
mechanisms at specific test conditions and appropriate models to assess it. Currently, direct
measurements of the interface toughness in complete TBC systems are still sparse because of
the complex shape and the anisotropy of the TBC systems. Further studies are needed to
compensate the current data base on this important issue.
2.7 Summary
In this chapter, the components of the multi-layered and multifunctional TBC system have been
reviewed regarding material requirements, microstructure and properties. The failure
mechanisms and the interface toughness measurement of different TBC systems were also
discussed. The establishment of reliable lifetime models of TBCs requires a more
comprehensive understanding of the failure mechanisms of TBCs that involve material
structures and properties. In addition, further research and development (R&D) including new
coating design to improve the durability of TBCs is still crucial to meet the growing demands
in the gas turbine engine industry.
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
71
Chapter 3 Pt Effect on Early Stage Oxidation Behaviour of Pt-
diffused γ-Ni/γ’-Ni3Al Coatings
3.1 Introduction
The Ni-Al-based coating, which is commonly used for oxidation protection of the Ni-base
superalloys, is designed to develop a protective Al2O3 (thermally grown oxide, TGO) layer
during high temperature exposure. This coating fails when the TGO spalls off, which is driven
by the increasing TGO strain energy as the oxide thickens. Thus, the lifetime of the coating
usually depends on the spallation resistance of the TGO layer. Numerous studies have reported
that the addition of platinum (Pt) can significantly improve TGO spallation resistance of Ni-
Al-based alloys or coatings for long-term oxidation at elevated temperatures. For instance, Y.
Chen et al. [159] concluded that Pt addition can improve the oxide spallation resistance of the
γ/γ’ nickel aluminide alloys mainly due to the selective oxidation of aluminium promoted by
Pt. P. Y. Hou and V. K. Tolpygo [44] found that 5 - 8 at.% Pt can significantly enhance the
oxide spallation resistance of nickel aluminide coatings during cyclic oxidation. They pointed
out that Pt can prevent the segregation of impurities (e.g. sulfur) at the oxide/coating interface,
resulting in enhanced interfacial adhesion and improved TGO lifetime. Other researchers [45,
48, 68, 160] have reported this benefit of Pt caused by other mechanisms such as inhibit ing
void formation at the coating/oxide interface.
It is widely accepted that for Al2O3-scale forming Ni-Al-based coatings, one (usually θ) or
more transient forms of Al2O3 initially form during the early stage oxidation, followed by the
eventual transformation to stable α-Al2O3 [89, 91, 161]. Previous studies have suggested that
the early stage oxidation behaviour can affect the prolonged oxidation behaviour of the Ni-Al
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
72
coating. For example, V. K. Tolpygo and co-workers [162] found that different transient
alumina to stable alumina transition rate during early stage oxidation leads to local TGO
thickness heterogeneity, which can induce local tensile stresses and is detrimental to TGO
lifetime for prolonged oxidation. Therefore, it is important to investigate the early stage
oxidation behaviour for a reliable estimation of TGO performance. However, despite the large
number of studies about Pt effect on long-term oxidation as mentioned above, only limited
attempts [53, 94, 98, 163] have focused on the impact of Pt on early stage of oxidation
behaviour of Ni-Al alloys or coatings. In addition, some controversial views regarding the Pt
effect on the early stage θ-Al2O3 to α-Al2O3 transformation have been reported. For instance,
J. Jedlinski et al. [94] found that Pt resulted in an earlier development of α-Al2O3 (faster θ- to
α-alumina transition) for a β-NiPtAl alloy at 1100°C oxidation, whereas Y. Cadoret, et al. [98]
concluded that Pt can slow down the θ-Al2O3 to α-Al2O3 transition by studying the early stage
oxidation of Ni50Al50 and Ni40Pt10Al50 alloys at 900 and 1100°C, respectively. H. Svensson and
co-workers [163] concluded that Pt has no effect on this phase transformation by studying the
initial oxidation of β-NiPtAl alloys. However, for the Pt-diffused γ-Ni/γ’-Ni3Al coating system,
little work has been done on the Pt effect on the θ-Al2O3 to α-Al2O3 transformation. The only
related work to the best of our knowledge, suggested that a higher Pt content led to a larger
amount of θ-Al2O3 in the oxide scale by photo-luminescence piezo-spectroscopy (PLPS) [131].
This indicates a slower θ-Al2O3 to α-Al2O3 transformation because of Pt addition for this γ/γ’
coating system. However, no mechanism has been proposed in their work about this Pt effect
on the phase transformation rate, mainly due to the lack of microscopic information about the
distribution of the two alumina polytypes and their relation to the coating composition (Pt
contents).
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
73
In this chapter, the early stage oxidation behaviours of γ/γ’ Ni-Al coatings with different Pt
additions are studied and compared in terms of TGO microstructure, transient oxide to stable
α-Al2O3 transformation rate and scale stress evolution, to provide a more comprehensive
understanding of the Pt effect on the early stage oxidation of this coating. Moreover, in order
to explore the mechanism of the Pt effect on alumina phase transition, the phase mapping of
the alumina polytypes (transient θ-Al2O3 and stable α-Al2O3) in the scale during the early stage
was conducted by the ASTAR (NanoMEGAS) automated crystal orientation mapping on
transmission electron microscopy (TEM) [164]. This technique offers significantly improved
lateral spatial resolution over the recently developed transmission electron backscatter
diffraction (t-EBSD) [165], which enables the diffraction pattern indexed phase mapping of
both θ-Al2O3 and α-Al2O3 during early stage oxidation for the first time, to the best of our
knowledge. A new transformation mechanism has been proposed based on these results to
explain the different transformation rate observed on coatings with different Pt contents.
3.2 Experimental procedures
3.2.1 Sample preparation and thermal treatment
CMSX-4 single crystal Ni-based superalloy (Table 3.1, Rolls-Royce plc) was used as substrates
throughout this study. The as-received superalloy bars were cut into buttons (20 mm diameter
and 3 mm height) using a SiC cutting blade in a precision cut-off machine (Accutom 5, Struers).
All substrates were ground and polished to 1 µm finish and then washed in acetone before Pt
electroplating. Platinum was electroplated on the substrate surface by using Q salt
(Tetraammineplatinum(II) hydrogen phosphate, Johnson Matthey) (See Table 3.2). The anode
was the platinum foil and the cathode was the CMSX-4 substrate. The electroplating
temperature was maintained in the required range (91~93 °C) for sufficient cathode efficiency.
The pH value was maintained in the recommended pH range (10.0-10.6) by regular addition
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
74
of ammonia solution during the electroplating process. In addition, the regular addition of
distilled water was necessary in order to maintain the volume of the bath. A magnetic stirrer
was applied in the bath to maintain a uniform Pt concentration of the bath. Three groups of
samples (each group contains three button samples) were fabricated, with electroplating time
0 min, 20 min and 50 min, respectively. All electroplated samples were washed in hot distilled
water (~80 ℃) for 0.5 h to remove the remaining salts from the electroplating bath. The
electroplated Pt layer thickness of each sample was examined by the focus ion beam system
(FIB; FEI Quanta 3D). Finally, the substrates with different electroplating time were annealed
in vacuum at 1150 °C for 2 h to obtain γ/γ’coatings with different Pt contents.
Table 3-1 Composition of CMSX-4 substrates
Element Ni Al Cr Co Ta Ti W Re
Wt. % 61.4 5.6 6.4 9.6 6.6 1.0 6.4 2.9
Table 3-2 Electroplating platinum bath
Chemical formula [Pt(NH3)4](HPO4)
Platinum content
pH
Temperature
Cathode current density
4~6 g/L
10.0~10.6 (optimum 10.5)
91~93 °C
2~7 mA/cm2
Isothermal oxidation of the γ/γ’ coatings with different Pt contents was performed at 1000ºC
in an automated rig (CMTM) in laboratory air. The samples were oxidized for different periods
of time up to 30 min, followed by air quenching.
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3.2.2 Luminescence measurement and data processing
Phase identifications of alumina polytypes as well as measuring residual stresses in oxide scales
was carried out at room temperature using PLPS on a Renishaw Invia Raman system
(RenishawTM, Gloucestershire, UK) with an argon laser source (λ=633 nm). The laser spot size
was ~5 µm. Although according to the subsequent results, some of the oxide scales studied
here have a Ni-oxide over-layer above the alumina layer, it was shown by Lipkin and Clarke
[83] that this over-layer does not prevent either the excitation or collection of the fluorescence
signal from the Cr3+ in alumina scale.
Before each experiment, the spectrometer was calibrated by taking a spectrum from a standard
pure silicon sample. Cr3+ fluorescence spectra were collected for each measurement with one
second acquisition time. Measurements for each sample were taken on a square grid of 200×
200 µm with a pitch of 20 µm, thus one map contains 121 measurement points for each sample.
To determine the peak positions in each spectrum, all spectra were deconvoluted in Wire 4.2
software (RenishawTM) with an automatic fitting program by two mixed Gaussian-Lorentzian
functions [166]. It is well accepted that the R1 and R2 doublet at ~14400 cm-1 is produced by
Cr3+ fluorescence in α-Al2O3, and θ-Al2O3 produces luminescence peaks at ~14546 cm-1 (T1),
~14626 cm-1 (T2) and ~14330 cm-1 (T3) [86], as shown in Fig. 3.1 a. The relative intensities
of the α-Al2O3 and θ-Al2O3 luminescence lines provide a semi-quantitative indicator of the θ-
Al2O3 content in the spot area, which is given by [86]:
Cθ=A𝑇1+𝐴𝑇2+𝐴𝑇3
A𝑇1+𝐴𝑇2+𝐴𝑇3+𝐴𝑅1+𝐴𝑅2 (3.1)
where A𝑇1, 𝐴𝑇2 and 𝐴𝑇3 are peak areas for three characteristic peaks of θ-Al2O3, respectively,
and 𝐴𝑅1 and 𝐴𝑅2 are peak areas of the characteristic R1 and R2 peaks.
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Since the R2 line has a nearly linear dependence on stress, the stress calculation was based on
the peak shift of the R2 line with respect to that of the unstrained single crystal sapphire (Fig.
3.1 b) by [84]:
∆𝜈 = 5.07(𝑐𝑚−1𝐺𝑃𝑎−1)σ (3.2)
where ∆𝜈 is the frequency shift of the R2 line and σ is the residual stress by assuming an
equibiaxial plane stress state in the scale.
Fig. 3.1 a) Luminescence spectrum showing characteristic peaks for θ-Al2O3 and α-Al2O3; b)
R peaks of an α-Al2O3 scale on a Pt-diffused γ/γ’ coating after isothermal oxidation at 1100°C
for 1 h. For comparison, the red line shows the spectrum of a stress-free polycrystalline alumina.
3.2.3 ASTAR automated crystal orientation mapping on TEM
The high resolution phase mapping of the alumina polytypes was performed using the ASTAR
automated crystal orientation mapping on TEM (transmission electron microscope, FEI Tecnai,
F30) [167], operating at 300 keV. FIB (FEI Quanta 3D) in-situ lift-out technique [168] was
used to prepare site-specific specimens for this analysis. Firstly, the electron beam scanned
over the area of interest on the FIB sample, and the digital image of diffraction patterns were
saved by a high-rate digital camera at each point with a spot size of 5 nm. The crystal
information for each phase present in the scanned area was extracted from crystal information
files (CIFs) produced from powder diffraction data [169-171]. Theoretical generated patterns
b) a)
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(the templates) for each phase were then produced using the pattern generation method by
Zaefferer [172]. The acquired experimental patterns were automatically matched with the
generated templates by the so-called template matching process [173]. The degree of matching
is quantitatively given by the correlation index Q, given by:
𝑄 =∑ 𝑃(𝑥𝑖,𝑦𝑖)𝑇(𝑥𝑖,𝑦𝑖 )𝑚𝑖=1
√∑ 𝑃2(𝑥𝑖,𝑦𝑖)𝑚𝑖=1 √∑ 𝑇2(𝑥𝑖,𝑦𝑖 )
𝑚𝑖=1
(3.3)
where P(x, y) is the intensity function of a pattern; T(x, y) is the intensity function of the
corresponding template; x and y are bounded by the picture width and height respectively.
Furthermore the reliability index assesses the likelihood of the match being unique and is
calculated by [173]:
R=100(1-Q2/Q1) (3.4)
where Q1 and Q2 stand for the two highest values of the correlation indexes for one pattern.
For a practical point of view, a reliability >15 is considered to be a safe solution.
3.2.4 Other characterization methods
An optical profilometer (ContourGT, Bruker) was used to examine the surface roughness of
the as-fabricated coatings. Scanning electron microscopy (SEM, FEI Quanta 650) and an
optical microscope (Olympus BH2-UMA) were used to examine the coating surface
microstructure after oxidation. Investigations of cross-sectional microstructure of oxide scales
were conducted by FIB (FEI Quanta 3D). Conventional angular dark field (ADF) images were
obtained by TEM (FEI Tecnai, F30) with STEM (scanning transmission electron microscope)
detector on the same samples for the orientation mapping in Section 3.2.3.
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3.3 Results
3.3.1 Microstructure of the as-fabricated coatings with different Pt contents
Table 3.3 shows the electroplating time, corresponding Pt layer thickness, average Pt
concentration of as-diffused coatings and root mean square roughness (Rq) of the coating
surface. 20 min and 50 min electroplating time result in 2 µm and 5 µm Pt, and the
corresponding average Pt concentration of the coating is 12.6 at.% and 21.3 at.%, respectively.
The root mean square roughness values are almost identical for the two as-diffused coatings,
and also very close to that of the uncoated substrate (~ 0.45 µm). The nearly identical roughness
values of all samples are prerequisite for the subsequent phase transformation rate analysis,
since the surface roughness has a significant effect on the alumina phase transformation in
oxide scales [162].
Fig. 3.2 shows the cross-sectional SEM images of the as-fabricated γ/γ’ coatings with 2 µm
and 5 µm Pt respectively. The bright phases of the coating are γ’ and the dark phases are γ. The
coating thickness and the fraction of γ’ phase in the coating both increase with increasing Pt
contents.
Table 3-3 Electroplating time and corresponding Pt thickness, average Pt concentration and
surface roughness
Electroplating time, min 0 20 50
Pt layer thickness, μm 0 2 5
Average Pt content, at. % 0 12.6 21.3
Rq, μm 0.45 0.40 0.48
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Fig. 3.2 Cross-sectional SEM images of as-diffused samples: a) 2 µm Pt coating; b) 5 µm Pt
coating.
3.3.2 Transient alumina to stable α-Al2O3 transformation
The alumina phase identification of scales on samples with different Pt additions after short-
term oxidation was carried out by PLPS. The spectra taken from different places on each
sample were highly reproducible. Therefore, only representative spectra are presented for each
sample. Fig. 3.3 and 3.4 show the PLPS spectra for all samples after 2 min and 10 min oxidation
at 1000 ºC respectively. The characteristic peaks for θ-alumina (if present) and α-alumina were
marked in each spectrum, and the semi-quantitative fractions of θ-alumina in the oxide scale
were calculated according to equation (1) based on the average of 10 randomly chosen spectra
for each sample.
It is clear that for no Pt addition sample, only stable α-Al2O3 can be identified without any
signals from θ-Al2O3 after just 2 min oxidation at 1000 ºC (Fig. 3.3 a). In other words, θ to α-
Al2O3 transition finishes within 2 min for samples without Pt. Conversely, a large amount
(~55.8 %) of θ-alumina is detected in the scale on 5 µm Pt coating after 2 min oxidation (Fig.
3.3 c), and for the coating with 2 µm Pt, nearly half of the scale consists of the θ polymorph
(Fig. 3.3 b). After 10 min oxidation, α-Al2O3 gradually prevails in the scale of 2 µm Pt coating
5 µm 10 µm
γ
γ’
~ 10 µm coating thickness ~ 20 µm coating thickness
γ’ γ
a) b)
2 µm Pt 5 µm Pt
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but a weak photoluminescence signal from θ-Al2O3 is still detected (Fig. 3.4 b). And the scale
of 5 µm Pt coating still contains a high content (~ 43.9 %, Fig. 3.4 c) of θ-Al2O3 after 10 min.
The θ-Al2O3 fractions of all samples during early stage oxidation are summarized in Fig. 3.5.
As can be seen, after 30 min oxidation, the θ-Al2O3 to α-Al2O3 transformation almost finish on
2 µm Pt sample (only ~ 5.0 % θ-Al2O3), while on 5 µm Pt coating, ~ 17 % θ-Al2O3 was still
detected. Clearly, the fraction of θ-Al2O3 decreases as the oxidation time increasing during
early stage oxidation of the two coatings with Pt addition. Moreover, the more Pt addition
results in more θ-Al2O3 in the scale, which supports the results in [98, 99, 131]. Hence, it can
be concluded that the transient θ-Al2O3 to stable α-Al2O3 transformation is retarded due to Pt
addition for the Pt-diffused γ/γ’ coating.
Fig. 3.3 Typical luminescence spectrum of a) 0 Pt, b) 2 µm Pt and c) 5 µm Pt samples after 2
min oxidation at 1000 ºC.
b)
c)
a)
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Fig. 3.4 Luminescence spectrum of a) 0 Pt, b) 2 µm Pt and c) 5 µm Pt samples after 10 min
oxidation at 1000 ºC.
Fig. 3.5 Fraction profiles of θ-Al2O3 as a function of oxidation time.
a)
b)
c)
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3.3.3 TGO composition and microstructure evolution
Fig. 3.6 shows the FIB cross sections of three samples after short-term oxidation at 1000ºC.
After 2 min, it can be seen from Fig. 3.6 a - c that the oxide scale thickness increases as the Pt
content increases. This is due to the fact that Pt addition can promote the growth of transient θ-
Al2O3, as has been confirmed in section 3.3.2, which grows much faster than the stable α-Al2O3.
It is worth noting that at this stage, all scales are composed of pure alumina. After 10 min
oxidation, however, the oxide scales of no Pt and 2 µm Pt sample become duplex in structure:
with outer transient oxide (NiO and/or spinel) layer and inner alumina layer (Fig. 3.6 d and e).
As for 5 µm Pt coating with the highest Pt content, the scale still composes of pure alumina
after 10 min oxidation, without any Ni-oxides formation (Fig. 3.6 f).
Fig. 3.6 a) - c): FIB/SEM cross-sectional images after 2 min oxidation; d) - f): FIB/SEM cross-
sectional images after 10 min oxidation at 1000 ºC of no Pt sample, 2 µm Pt coating and 5 µm
Pt coating, respectively.
Fig. 3.7 shows surface images and FIB cross-sectional images of three samples after 30 min
oxidation. The surface of no Pt sample exhibits fine faceted morphology consisting of NiO
grains (Fig. 7 a), and larger faceted spinel grains are observed on the 2 µm Pt coating (Fig. 3.7
b). While the surface of 5 µm Pt coating shows a combination of nodules and short whiskers
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(Fig. 3.7 c). The short whisker morphology is typical for an oxide growth as a result of outward
cation diffusion, as expected for the growth of the transient θ-Al2O3 at the gas/oxide interface
in this temperature regime [91, 98]. The short whiskers are only observed on the surface of 5
µm Pt coating. This is in agreement with the results in Fig. 3.5, which shows that after 30 min
oxidation, θ-Al2O3 to α-Al2O3 transition has almost completed on the no Pt sample and 2 µm
Pt coating. As for the cross-sectional microstructure, the oxide scale formed on no Pt sample
is multi-layered, consisting of an outer NiO layer, an intermediate spinel layer and an inner
alumina layer (Fig. 3.7 d). The oxide scale on 2 µm Pt coating is similar to that on no Pt sample,
but without the outermost NiO layer (Fig. 3.7 e). The scales on these two samples are not fully
dense, and a large number of pores can be observed at the interface of spinel/alumina. The
formation of the pores is attributable to the solid-state reaction between NiO and Al2O3, which
induces volume contractions and leads to formation of pores [30]. On the other hand, a dense
and uniform alumina layer (except for some Cr2O3 particles as shown by the arrow in Fig. 3.7
f) is observed on 5 µm Pt coating. After 30 min oxidation, the oxide scale of no Pt and 2 µm
Pt coating is much thicker than that of 5 µm Pt coating because the growth of Ni-oxides is
much faster than alumina.
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Fig. 3.7 a) - c): SEM surface images; d) - f): FIB/SEM cross-sectional images of no Pt sample,
2 µm Pt coating and 5 µm Pt coating after 30 min oxidation at 1000 ºC.
As shown in Fig. 3.6 and Fig. 3.7, sufficient Pt addition results in an exclusive alumina TGO
layer, compared to samples without Pt (multi-layer TGO structure) or with less Pt addition
(duplex TGO structure). This confirms that Pt can promote alumina growth by inhibiting Ni-
oxides growth at very early stage of oxidation.
3.3.4 TGO growth rate & stress evolution
Fig. 3.8 is the oxidation time at 1000 ºC vs. TGO scale thickness plot of the sample without Pt
and with 5 µm Pt, respectively. As can be seen, initially, 5 µm coatings exhibit a faster TGO
growth compared to the no Pt sample. This is because within 2 min oxidation, θ-Al2O3 to α-
Al2O3 transition finishes in the scale on no Pt sample (Fig. 3.3 a) and the TGO only composes
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of the slow-growing α-Al2O3. While fast-growing θ-Al2O3 accounts for a large proportion of
the scale on 5 µm Pt-diffused γ/γ’ coating (Fig. 3.3 c), which grows much faster.
Fig. 3.8 Oxide scale thickness evolution of no Pt sample and 5 µm Pt sample as a function of
oxidation time.
However, after 30 min oxidation, the average scale thickness of 5 µm Pt sample is only half of
that on no Pt sample (Fig. 3.8). The TGO growth rate of 5 µm Pt sample is significantly reduced
due to the phase transformation θ- to α-Al2O3, and α-Al2O3 gradually becomes dominant in the
scale of this sample without any Ni-oxide layer formation (Fig. 3.5 and Fig. 3.7 f). But for no
Pt sample, the oxide growth rate increases significantly as a result of Ni-oxide layers formation
(Fig. 3.7 d). Therefore, during the early stage of oxidation, Pt firstly increases oxide growth by
promoting growth of θ-Al2O3, and then it significantly slows down the oxide growth by
promoting the selective oxidation of Al and inhibiting the formation of Ni-oxides/spinel. In
other words, our study confirmed that Pt has two effects on the oxide scale composition during
the early stage oxidation: 1) it can slow down θ- to α-Al2O3 phase transformation; 2) Pt can
promote the selective oxidation of Al at very early stage of oxidation, in agreement with [53].
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The TGO stress evolution of samples without Pt and with the highest Pt content (5 µm) as a
function of oxidation time is shown in Fig. 3.9. After 2 min oxidation, the TGO stresses of the
two samples are similar. However, afterwards, the compressive TGO stress of no Pt sample
increases significantly from ~ - 2.4 GPa (2 min) to ~ - 3.8 GPa (30 min), whereas for the 5 µm
Pt sample the compressive stress gradually decreases from ~ - 2.2 GPa (2 min) to ~ - 1.4 GPa
(30 min). This simply implies that Pt addition can reduce the compressive TGO stress during
the early stage oxidation of γ/γ’ Ni-Al coatings.
Fig. 3.9 The TGO stress evolution of no Pt sample and 5 µm Pt sample as a function of
oxidation time.
3.3.5 Automated crystal orientation mapping with TEM
Fig. 3.10 and Fig. 3.11 show the results of conventional ADF-STEM and crystal orientation
mapping on TEM of samples with 2 μm Pt coating and 5 μm Pt coating respectively after 10
min oxidation at 1050 ºC. ADF-STEM allows the overall structure (coating-alumina-plat inum
capping layer) to be determined. However, the similarity in composition of the different
alumina polytypes in the scale means that these cannot be determined by the ADF-STEM
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contrast. Instead scanning diffraction analysis was used to identify the polytypes and the crystal
orientation of the alumina at a high spatial resolution (~ 5 nm). TEM diffraction patterns were
recorded in a raster scan across the sample and then compared to libraries of templates
generated for the different crystal structures present in the sample.
Fig. 3.10 a) ADF STEM image of 2 µm Pt sample after oxidation at 1050 ºC for 10 min; b)
combined phase map and phase reliability map obtained from automated crystal orientation
mapping in TEM, taken from the red box region in a). Green: θ-Al2O3; red colour: α-Al2O3.
From the analysis, maps of the θ-Al2O3 (indicated by green pixels in Fig. 3.10 and 3.11) and α-
Al2O3 (indicated by red pixels) could be obtained. These phase maps were overlaid on phase
reliability maps (see Eq. 4) to allow grain details of the underlying coating to be shown. The
general effect of platinum slowing down θ to α transformation appears to be borne out by this
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analysis, by inspecting the proportion of α-Al2O3, which was higher in Figure 3.10 b compared
to Figure 3.11 b. This is in agreement with the results in Section 3.3.2.
Fig. 3.11 a) ADF STEM image of 5 µm Pt sample after oxidation at 1050 ºC for 10 min; b)
combined phase map and phase reliability taken from the red box region in a). Green colour:
θ-Al2O3; red: α-Al2O3.
The spatial distribution of θ and α-Al2O3 in the scale can be seen in the maps. Generally, both
maps (Fig. 3.10 b and Fig. 3.11 b) reveal a fine level of detail with grains as small as ~10 nm
in the inner region of the scale, while large columnar grains (100 - 200 nm) typically exists in
the outer region of the scale. This is consistent with a previous study on NiAl alloy by bright
field TEM imaging [163]. Importantly there is variation in the distribution of the different
polytypes along the surface of the coating. By inspecting individual grains which can be
distinguished by the greyscale contrast in the phase reliability map (e. g. the regions ‘A’ and
‘B’ in Figure 3.10 b are distinct grains), there are regions where α-Al2O3 has formed at the
coating surface, while at others θ-Al2O3 is present. Looking further across the scale, it is clear
that some regions have not transformed from θ to α as there are columnar grains of θ-Al2O3
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extending across the entire scale thickness. These observations suggest that the θ-Al2O3 to α-
Al2O3 transformation rate is inhomogeneous along the coating/oxide interface and related to
the underlying coating grains.
3.3.6 PLPS studies on Ni-Al-Pt alloy samples
From PLPS macroscopic results (Section 3.3.2), we conclude that higher Pt content results in
slower θ-Al2O3 to α-Al2O3 transition. However, the mechanism of this Pt effect is still unclear.
Given the absence of any strong crystallographic coherency between the alumina grains (near
the coating surface) and the coating grains (Appendix. A), it is therefore important to
investigate the effect of the local platinum distribution in the coating since the inhomogeneous
distribution of platinum near the surface of the coating could result in the inhomogeneous
nucleation of α-Al2O3, as confirmed in Section 3.3.5.
To validate the above assumption, a direct method is to examine the coating grains individually
to determine their phase structure and correlate this to the oxide phase grown on each grain
during the early stage oxidation of Pt-diffused γ/γ’ coatings. Nevertheless, there are problems
with this method. Primarily, the alumina growth results in the depletion of aluminium in the
coating and hence the γ’ to γ phase transition can take place. Thus the TEM samples did not
retain any of the original γ/γ’ phase distribution even for very short oxidation exposure.
Furthermore, the optical microscopy was also unable to distinguish the underlying γ’ or γ phase
from the oxide surface of these coatings.
In order to circumvent these problems, the early stage oxidation behaviour of Ni-Al-Pt alloy
with γ/γ’ microstructure was studied. Y. Chen et al. [159] have shown that this alloy maintains
phase contrast of γ’ and γ phase under optical microscope on the oxide surface after up to 4 h
oxidation at 1150 ºC. In addition, γ’ to γ phase transformation is prevented during the early
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stage oxidation due to a sufficient aluminium source in the underlying alloy allowing the
ordered γ’ phase to exist even after aluminium depletion. Hence, it is possible to study the early
stage oxidation of this alloy by PLPS with relation to the underlying alloy phase structure to
verify the above assumption.
Ni-20Al-xPt (x= 0, 10, 20 at.%) alloys with γ/γ’ microstructure were investigated. The optical
surface images of as-received alloys are shown in Fig. 3.12. The bright areas in Fig. 3.12 a - c
represent γ’ and the dark dendritic areas consist of tiny γ’ blocks embedded in the γ network-
structure channels (Fig. 3.12 d). With increasing Pt content, the dendrites become narrower and
the phase fraction ratio of γ’ to γ increases.
The oxides grown on γ’ and γ/γ’ region, respectively, were studied by PLPS to analyse the
phase after short time exposure at 1050 ºC. The results of three alloys are shown in Fig. 3.13-
3.15, respectively. The γ’ and γ/γ’ region can be easily distinguished after oxidation for all
alloys. As the Pt content increases, the θ-Al2O3 fraction also rises, no matter which region (γ’
or γ/γ’) is detected. This also confirms Pt effect on retarding θ-Al2O3 to α-Al2O3 transformation.
For Ni-20Al alloys, the average θ-Al2O3 fraction of γ’ regions (5 measurements on random
locations) is ~ 21.8 % after 2 min oxidation (Fig. 3.13 b), while θ-Al2O3 fraction of γ/γ’ regions
at this stage is 14.9 % (Fig. 3.13 c), which is 31.6 % lower than that on γ’ areas. After 10 min,
the average θ-Al2O3 fraction of γ’ areas for Ni-20Al is ~ 9.2 % (Fig. 3.13 e), whereas very
weak θ-Al2O3 signal can be detected on γ/γ’ areas (Fig. 3.13 f). The results of Ni-20Al-10Pt
and Ni-20Al-20Pt alloys also show the same trend. For instance, after 2 min, θ-Al2O3 fraction
of γ/γ’ areas of Ni-20Al-10Pt (53.4 %, Fig. 3.14 c) is 30.9 % lower than that on γ’ areas (77.7 %,
Fig. 3.14 b). And for Ni-20Al-20Pt alloy after 10 min oxidation, θ-Al2O3 fraction of γ/γ’ areas
is 35.8 % lower than that on γ’ areas (Fig. 3.15 e and f). The results in Figs. 3.13 - 3.15 have
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been summarised in Table 3.4 and 3.5. From Table 3.4 and 3.5, it can be concluded that θ-
Al2O3 fraction of γ/γ’ areas is significantly lower than that on γ’ areas for three alloys
throughout the early stage oxidation, which coincides with our assumption that γ phase can
promote θ-Al2O3 to α-Al2O3 transformation, resulting in a lower θ-Al2O3 fraction of γ/γ’ regions.
In other words, these results support that γ grains near the coating/oxide interface promote the
transformation from θ-Al2O3 to α-Al2O3 while γ’ grains retard it.
Fig. 3.12 Microstructure of the as-received a) Ni-20Al, b) Ni-20Al-10Pt, c) and d) Ni-20Al-
20Pt alloy. The inset in d) shows the magnified morphology of the tiny γ channels in γ/γ’ region.
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Fig. 3.13 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al alloy after oxidation at 1050 ºC for 2 min and 10 min, respectively.
Fig. 3.14 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al-10Pt alloy after oxidation at 1050 ºC for 2 min and 10 min.
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Fig. 3.15 a) and d): the optical image; b) c) e) and f): PLPS spectra taken from different phase
regions of Ni-20Al-20Pt alloy after oxidation at 1050 ºC for 2 min and 10 min.
Table 3-4 θ-Al2O3 fraction of γ/γ’ areas and γ’ areas, respectively for the three Ni-Pt-Al
alloys after 2 min oxidation at 1050 ºC
Alloys Cθ (%), γ/γ’ areas Cθ (%), γ’ areas
Ni-20Al 14.9 21.8
Ni-20Al-10Pt 53.4 77.7
Ni-20Al-20Pt 73.2 100
Table 3-5 θ-Al2O3 fraction of γ/γ’ areas and γ’ areas, respectively for the three Ni-Pt-Al
alloys after 10 min oxidation at 1050 ºC
Alloys Cθ (%), γ/γ’ areas Cθ (%), γ’ areas
Ni-20Al 0 9.2
Ni-20Al-10Pt 21.6 46.2
Ni-20Al-20Pt 45.3 70.6
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3.4 Discussion
3.4.1 Pt effect on the θ-Al2O3 to α-Al2O3 transformation
Platinum addition has been shown to alter the transformation dynamics of θ-Al2O3 to α-Al2O3
[94, 98]. However, the precise mechanism for this is still unclear. The results in this article
suggest that one effect of platinum is the local variation in the nucleation of α-Al2O3 in the θ-
Al2O3 scale. This alters the conventional model of oxide transformation on these coatings.
According to Hayashi and Gleeson [93], α-Al2O3 nucleates uniformly along the coating/oxide
interface, leading to a duplex layered structure of the oxide scale, as shown schematically in
Fig. 3.16 a. However, in this study, γ grains near the interface promote α-Al2O3 nucleation
while γ’ grains retard this transformation (Fig. 3.16 b and c). Under this circumstance, we can
explain why Pt addition slow down this transformation and the non-uniform distribution of α-
Al2O3 along the interface. The θ-Al2O3 to α-Al2O3 transformation is accelerated on coatings
with lower Pt content due to more γ grains near the coating/oxide interface (Fig. 3.16 b).
Conversely, for coatings with higher Pt content, this transformation is retarded because of the
higher γ’ phase to γ phase fraction ratio near the coating/oxide interface (Fig. 3.16 c). Then,
the early α nuclei on γ grains grow toward the gas/oxide interface and finally grow into large
grains, while α-Al2O3 nucleation happens later on γ’ grains, which is consistent with the
microstructural observations in Fig. 3.10 b and Fig. 3.11 b. A previous study suggested that β
grains of NiCrAlY coatings have a large tendency to form θ-Al2O3 than the γ grains [174],
which seems to be supportive to our mechanism considering that both γ’ and β phase contain
more aluminium than the γ phase because aluminium diffusion might affect the θ to α
transformation rate [98].
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Fig. 3.16 a) Previously reported growth model of α-Al2O3: uniformly nucleate along the
interface, resulting in a layered structure of the oxide scale; b) and c) new growth model in our
study which illustrates lower Pt content and higher Pt content coatings, respectively.
Through the use of spatially resolved PLPS measurements on the NiPtAl alloys with γ/γ’
structure, it was possible to show that where Pt stabilises the ’ structure in nickel the
suppression of the θ-Al2O3 to α-Al2O3 transition was seen. This offers a validation for the
proposed new model in Fig. 3.16. In the absence of significant crystallographic coherence
between the alumina grains (both α and θ) and the coating (Appendix A), it is concluded that
the original distribution of ’ in the coating (associated with the local Pt content) directly
determines the phases present in the oxide scale. This study combined with other studies on
MCrAlY coatings [55, 174] can provide a thorough understanding of the relationship between
the oxide phase/composition and the underlying coating microstructure.
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In search of the reason(s) why the alumina phase transformation is retarded on the ’ phase
region, firstly, as a noble metal, Pt is non-reactive and does not form any oxide during oxidation
of the coatings. There is also no evidence that Pt can diffuse into the lattice of Al2O3 or
segregate to the grain boundaries of Al2O3 in the literature. Therefore, it is very unlikely that
the retarded θ to α transformation on the γ’ region is caused by the direct chemical interaction
between Pt and Al2O3.
On the other hand, since Pt strongly partitions into γ’, the γ and γ’ phase have different lattice
parameters and corresponding interplanar spacing (d-spacing) [159]. This would affect the
epitaxial strain in the θ-Al2O3 at the very initial stage of oxidation. A low d-spacing mismatch
between θ-Al2O3 and the underlying metal phase is expected to favour the formation of θ-Al2O3
due to the low energy barrier, thus slowing down the θ to α phase transformation. For this
reason, it is proposed that a possible reason for the retarded θ to α transformation on the γ’
regions is the lower minimum d-spacing mismatch between θ-Al2O3 and γ’ phase.
To support this argument, the possible orientation relationships in terms of minimum
interplanar spacing (d-spacing) mismatch between θ-Al2O3 and two coating phases (γ’ and γ)
were investigated. In order to find the possible orientation relationships between θ-Al2O3 and
original γ’ (and γ) phases, the transmission diffraction data including the d-spacing values for
γ’, γ and θ-Al2O3 phase were generated by SingleCrystal ™ software with the basic lattice
parameters for each phase as input [171, 217]. Note that the lattice parameters for γ’ and γ were
measured at 1000℃ in [217], while the lattice parameters for θ-Al2O3 was measured at room
temperatures [171]. Thus the d-spacing values of θ-Al2O3 has been multiplied by a factor of
(1+αAl2O3∆𝑇), where αAl2O3 is the coefficient of thermal expansion of alumina (~8 ppm℃-1) and
∆𝑇 is the temperature drop between the oxidation temperature and room temperature
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
97
(~1000℃). Table 3.6 gives the data for the two Ni-based phases (γ’ and γ) and θ-Al2O3
including the low-index (hkl) planes and the corresponding d-spacing values.
Table 3-6 Low-index (hkl) planes and the corresponding d-spacing values of two Ni phases
(γ’ and γ) and θ-Al2O3: the d-spacing mismatch between planes of θ-Al2O3 and the
corresponding γ’ (or γ) plane with closest d-spacing matching is calculated and listed as the
strain
θ-Al2O3 γ’ γ
(hkl) dhkl, Å (hkl) dhkl, Å strain (hkl) dhkl, Å strain
(201̅) 4.5307
(201) 3.5542
(400) 2.8660
(1̅11) 2.5700
(111) 2.4497
(310) 2.3172
(202) 2.2652
(311) 2.0298
(1̅12) 2.0232
(11̅2) 1.9078
(003) 1.8211
(2̅03̅) 1.6273
(1̅13) 1.5728
(113) 1.4909
(020) 1.4562
(220) 1.4113
(221) 1.3475
(222̅) 1.2850
(404) 1.1326
(223) 1.0854
(024) 0.9962
(100) 3.6000 1.30%
(110) 2.5456 0.96%
(111) 2.0785 -
(200) 1.8000 1.17%
(210) 1.6100 1.00%
(211) 1.4697 0.92%
(220) 1.2728 0.96%
(221) 1.2000 -
(310) 1.1384 0.51%
(311) 1.0853 0.009%
(222) 1.0329 -
d-spacing mismatch between
planes of θ-Al2O3 and γ’:
ε=𝑑ℎ𝑘𝑙𝜃 −𝑑ℎ′𝑘′𝑙′
𝛾′
𝑑ℎ𝑘𝑙𝛾′
For example, the closest d-
spacing mismatch between θ-
Al2O3 and γ’:
ε=𝑑(223)𝜃 −𝑑
(311)
𝛾′
𝑑(311)𝛾′ =0.009%
(100) 3.5623 0.23%
(110) 2.5189 2.03%
(111) 2.0567 -
(200) 1.7725 2.74%
(210) 1.5931 2.14%
(211) 1.4543 0.13%
(220) 1.2533 2.53%
(221) 1.1874 -
(310) 1.1265 0.54%
(311) 1.0689 1.54%
(222) 1.0234 -
d-spacing mismatch between
planes of θ-Al2O3 and γ:
ε=𝑑ℎ𝑘𝑙𝜃 −𝑑ℎ′𝑘′𝑙′
𝛾
𝑑ℎ𝑘𝑙𝛾
For example, the closest d-
spacing mismatch between θ-
Al2O3 and γ:
ε=𝑑(223)𝜃 −𝑑
(311)
𝛾
𝑑(311)𝛾 =1.54%
*Note: The data for θ-Al2O3 is listed as a decreasing order for d-spacing values, and only the
d-spacing values which are within the range of γ’ and γ phases are listed.
The d-spacing mismatch between planes of θ-Al2O3 and the corresponding γ’ (or γ) plane with
closest d-spacing matching is defined as:
ε=𝑑ℎ𝑘𝑙𝜃 −𝑑ℎ′𝑘′𝑙′
𝛾′
𝑑ℎ𝑘𝑙𝛾′
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
98
As can be seen from Table 3.6, in the case of 11 low-index planes of γ’ (or γ), there is a clear
d-spacing mismatch advantage which favours the coherent growth of θ-Al2O3 on the γ’ region
by working through the γ’ (or γ) planes as follows:
3 of the 11 low-index planes do not have a close match to the planes of θ-Al2O3 (these planes
are (111), (221) and (222)). As for the remaining 8 planes, 6 of them have a d-spacing mismatch
(strain, ε) below 1% between θ-Al2O3 and γ’ (with the other two exhibiting a strain up to 1.3%).
On the contrary, for the γ planes, only three of them have a strain below 1% and the others
show a much higher strain (2 - 2.5%) than that for the γ’ phase. 1% has been chosen as a
reasonable threshold for the elastic limit of alumina beyond which the coherency is likely to
fail. Hence, there are many more possible coherent orientations between θ-Al2O3 and γ’ for a
given surface facet of the metal substrate. This accounts for the strong tendency for the growth
of θ-Al2O3 on the γ’ phase region, rather than the γ phase.
3.4.2 Pt effect on TGO composition & stress
The benefit of Pt for promoting exclusive alumina growth and inhibiting Ni-oxide growth, as
already discussed in previous publications [53, 68, 159, 175], can be ascribed to two reasons.
Firstly, Pt is inert and does not react with oxygen at temperatures up to 1200 ºC. But it has
strong preference of replacing Ni site in the ordered L12 structure of γ’, resulting in an increase
of the Al:Ni atom ratio with increasing Pt content on a given crystallographic plane containing
both Al and Ni. This increment of Al:Ni ratio would kinetically favour the formation of alumina
in preference to NiO on that plane, when considering only the reacting constituents, Al and Ni
[53]. Secondly, the enrichment of Pt in the subsurface of the Pt-containing coatings during the
early stage oxidation also contributes to this Pt benefit [49]. This enrichment not only favours
the formation of alumina as stated in the first aspect, it can also decrease the chemical activity
of Al (aAl) at the oxide/coating interface. The decrease of aAl would, in turn, facilitate the
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99
diffusion flux of Al from the inner part of the coating to the oxide/coating interface [175], thus
providing a sufficient and consistent Al source for alumina growth.
Reduction of the TGO compressive stress of Pt-containing coatings during early stage
oxidation can be explained by the slower transformation from transient alumina to stable α-
Al2O3. Firstly, the slower transformation rate due to presence of Pt allows for an extended stress
relaxation process, which is due to the tensile stresses generated by the large volume shrinkage
(~10 %) associated with the transformation from the monoclinic θ phase to the hexagonal α
phase with higher density [162, 176]. This tensile stress relaxation can lead to significant
reductions of TGO compressive stress [177, 178]. Secondly, the increase ratio of θ/α in the
oxide scale in the presence of Pt may lower the growth stress since θ-Al2O3 is more plastic than
α-Al2O3, which also contributes to the lower stress of higher Pt coatings [98].
3.4.3 Early stage oxidation effect on prolonged oxidation performance
As mentioned above, Pt additions can significantly promote the growth of metastable θ-Al2O3,
which is less protective than the stable α-Al2O3. Moreover, the growth rate of θ-Al2O3 is about
an order of magnitude higher than α-Al2O3 [88]. An extended lifetime of less protective and
faster growing θ-Al2O3 due to Pt addition seems disadvantageous to the scale spallation
resistance at first sight. However, the long term oxidation investigations of coatings used in
this study (Appendix B) shows that coatings with highest Pt contents still presented the longest
TGO lifetime, despite the extended lifetime of the θ-Al2O3 scale. This is mainly attributed to
the fact that Pt can inhibit detrimental Ni-oxide formation at early stage of oxidation (Fig. 3.7).
In addition to the TGO lifetime, the early stage oxidation can also affect the scale morphology
and stress evolution during the prolonged oxidation. It is found that the slower θ to α transition
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
100
rate during early stage oxidation caused by Pt addition can result in a different oxide
morphology and low compressive stress of the stable scale, compared to coatings without Pt
(Appendix C). This stress reduction can also contribute to the longer TGO lifetime of Pt-
containing coatings, as reported elsewhere [98].
3.5 Summary
Pt addition has three effects on the early stage oxidation of γ/γ’-based Ni-Al coatings: 1) retard
θ-Al2O3 to α-Al2O3 transformation thus extending transient θ-Al2O3 lifetime; 2) promote the
growth of alumina and inhibit the growth of Ni-oxides; 3) significantly reduce the TGO stress
during the early stage of oxidation. Crystal orientation mapping results show that the nucleation
of α-Al2O3 is inhomogeneous along the oxide/coating interface and might be related to the
variation of coating compositions due to Pt additions. Spatially resolved PLPS study of Ni-Pt-
Al alloy (with γ/γ’ microstructure) shows that where Pt stabilises the γ’ structure in nickel, the
suppression of θ-Al2O3 to α-Al2O3 transition is observed. Based on these findings, a new
mechanism has been proposed to explain this Pt effect on θ-Al2O3 to α-Al2O3 transformation:
γ grains near the coating/oxide interface promote α-Al2O3 nucleation while γ’ grains retard this
transformation.
Appendix A. Coating/alumina orientation analysis
Three regions (red boxes in Fig. A1) were investigated to explore the possibility of coherent
relationships affecting the phase transformation in the TGO scale. Through a non-negative
matrix factorisation (NMF) decomposition [179], representative diffraction patterns for
different phases were isolated at the interface region, which also gave a general idea of their
location in the sample. Therefore, phase maps that were complementary to the ASTAR analysis
can be built and compared. Fig. A2 - A4 exhibit the results for Region 1 - 3.
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
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Fig. A1 Virtual bright-field image of the FIB-TEM sample (5 µm Pt coating after 10 min
oxidation at 1050 ºC).
For Region 1, a small region of θ-Al2O3 close to the nickel-alumina interface (the upper phase
map in Fig. A2) can be seen. The diffraction patterns for these θ-Al2O3 have partial coherency
with the underlying nickel (overlapping reflections as shown by the upper diffraction pattern
image). Indexing the remainder of the region, it was found that the majority of the alumina has
transformed to α-Al2O3 (lower phase map in Fig. A2). By overlaying these diffraction patterns,
both θ-Al2O3 and α-Al2O3 in this region have no strong orientation relationship with the nickel
substrate, as shown by the lower diffraction pattern image.
There is a twinned area in the middle of the nickel in Region 2 (Fig. A3, twin boundaries are
indicated). Similar to Region 1, a small coherent θ-Al2O3 region near the surface of nickel is
observed. In addition, the coherency is more prominent in the twinned region than the outer
twin region (as the greater coincidence in the diffraction pattern from the inner-twin suggests).
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Importantly, both α grains and large θ grains that form further out from the interface layer
exhibit little or no coherency with the substrate, as indicated by the lower diffraction patterns
in Fig. A3.
Fig. A2 Phase maps and corresponding diffraction pattern images of different phases in
Region 1.
Region 3 is predominantly α-Al2O3 (Fig. A4). Like other regions, Region 3 also shows no
orientation relationship between the alumina grains and the nickel substrate apart from a
reasonably coherent θ-Al2O3 orientation that is present in the lower part of the region (upper
diffraction pattern image).
Overall, θ-Al2O3 always nucleates partial-coherently along the nickel surface whereas with
continued growth this transforms into incoherent α-Al2O3, with no obvious relationship to the
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103
underlying nickel orientation. This suggests that the nickel substrate orientation has no effect
on θ to α-Al2O3 phase transformation.
Fig. A3 Phase map and corresponding diffraction pattern images in Region 2.
Fig. A4 Phase map and corresponding diffraction pattern images in Region 3.
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
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Appendix B. Prolonged oxidation lifetime of coatings with different
Pt additions
Fig. B1 shows the optical images of three coatings with different Pt additions after prolonged
oxidation. For no Pt sample, local TGO spallation occurs after 50 h (bright areas in Fig. B1 a),
and spallation is observed nearly all over the surface after 200 h (Fig. B1 c). In contrast, the 5
µm Pt coating scale remains intact up to 200 h oxidation (Fig. B1 e).
Fig. B1 Optical images of sample surface after different periods of prolonged oxidation at
1050 ºC. The bright area is the surface of the underlying metal due to spallation of the TGO.
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
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Appendix C. Early stage oxidation effect on the stable scale
morphology & stress
The TGO stress evolution of no Pt sample and 5 µm Pt coating after prolonged oxidation is
shown in Fig. C1. The TGO stress of 5 µm Pt coating is lower than that of no Pt sample, until
the TGO spallation on Pt sample which causes a significant drop of stress (Fig. C1 a). Moreover,
there are two stages of stress evolution of 5 µm Pt coatings. During Stage Ⅰ (up to 10 h), the
TGO stress remains at a low level (~ 3.5 GPa) compared to no Pt coatings. After 10 h (Stage
Ⅱ), the stress shows a significant rise (~ - 4.5 GPa). Due to Pt addition, the growth of θ-Al2O3
is extended and thus numerous θ-Al2O3 whiskers form on the surface of 5 µm Pt coating during
early stage oxidation. These whiskers totally transform into stable α-Al2O3 and sustain their
whisker shape after phase transformation on the surface of the stable scale (Fig. C1 b, after 5 h
oxidation). These whiskers, being less constrained by the underlying coating, can cause the R
peaks of α-Al2O3 in PLPS spectrum shift to much higher frequencies, resulting in lower TGO
stresses. In this case, it should be assured that the scale is intact, because any crack would give
similar stress-free signals, which has been ruled out by the microstructural study of the oxide
scale. While for no Pt sample, since θ to α transformation is much faster, no whiskers remain
on the stable scale. Therefore the TGO stress of this sample is higher until TGO spalls. It is
noticed that whiskers on the stable scale of 5 µm Pt coating become smoother with oxidation
time (Fig. C1 c) as a result of surface diffusion driven by a reduction of surface energy, which
causes a slight rise of stress during the Stage Ⅰ. Eventually, these whiskers on the surface
evaporate after prolonged exposure at high temperatures (Fig. C1 d), thus leading to a
significant rise of the TGO stress.
CHAPTER 3 PT EFFECT ON EARLY STAGE OXIDATION BEHAVIOUR OF γ/γ’ COATINGS
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Fig. C1 a) Residual stress evolution in TGO scale on no Pt sample and 5 µm Pt coating with
oxidation; b) - d) scale surface morphologies of 5 µm Pt coating after 5h, 10 h and 50 h
isothermal oxidation at 1050 ºC.
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
107
Chapter 4 Effect of Superalloy Substrate on the Lifetime and
Interfacial Toughness of Electron Beam Physical
Vapour Deposited Thermal Barrier Coatings
4.1 Introduction
Thermal barrier coatings (TBCs) are multilayered and multifunctional material assemblies that
have been widely applied for an improved performance and efficiency of gas turbine engines
[31]. State-of-the-art TBCs consist of a ceramic top coat of yttria stabilized zirconia (YSZ)
deposited by electron beam physical vapour deposition (EB-PVD) or plasma spraying, a β-
PtAl or γ/γ’ based diffusion or MCrAlY (M=Ni, Co or a mixture of two) overlay bond coat,
and a nickel-based superalloy substrate. Since the YSZ top coat is permeable to oxygen, the
bond coat oxidizes and forms a thermally grown oxide (TGO) scale (usually Al2O3) on top of
itself during service at high temperature. The TGO layer is subject to very large compressive
stress when the system is cooled due to the thermal mismatch with the metal substrate. Various
failure modes have been observed for TBCs, and the most common one for EB-PVD TBCs is
cracking at the bond coat/TGO interface leading to buckling delamination [110].
A number of studies have reported that the superalloy substrate composition can affect the
cyclic lifetime of TBCs [180-184]. For example, R. T. Wu et al. [185] found that the cyclic
lifetime varied significantly for TBCs with Pt-diffusion bond coats on different superalloy
substrates including SRR99, CMSX-4, etc., although the reason for this difference remains
unclear. B. A. Pint et al. [186] have pointed out that titanium (Ti) in the CMSX-4 superalloy
substrate was suspected to have degraded the alumina scale adhesion for Pt-diffusion bond
coats, thus was detrimental to the lifetime of TBC systems. However, they also found that the
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
108
poorer coating performance for CMSX-4 superalloy was not found in all cases because the Pt-
modified aluminide bond coat lifetime on CMSX-4 substrate was significantly higher
compared to other superalloys. Thus, more comprehensive analysis is needed to understand the
coating performance on CMSX-4 superalloy and the effect of Ti.
Another important factor affecting the TBC lifetime is the resistance to crack propagation
(toughness) of the relevant interfaces. The coating durability is dependent on the interplay
between the cracking driving forces and the resistance to crack propagation through the coating
or along the relevant interfaces. So measuring the relevant interface toughness of TBCs by
reliable test methods is critical to establish reliable lifetime models [149, 187]. Methods have
been developed to measure the interface fracture toughness in ceramic coating systems
including: (1) bending test of notched multilayer beams [134, 188]; (2) indentation [141, 189];
(3) blister methods [190]; (4) the push-out test [147, 149]. However, some of these methods
have limited applicability to TBC systems [132]. For instance, some TBCs on the turbine blade
samples cannot be tested by the methods which require specific geometry of samples, such as
push-out tests. Moreover, the crack path is difficult to control in a multilayer TBC system. In
some cases, cracks will divert away from the interface of interest into the vertical YSZ columns,
which makes it impossible to measure that interface. Recently, X. Wang et al. [132] have
utilized a cross-sectional indentation (CSI) method to measure the fracture energy of the EB-
PVD TBCs with the Pt-diffused bond coats. CSI does not require a designed sample geometry
and can produce controlled interface fracture. But complex finite element modelling is required
to analyse the plastic deformation of the substrate for the fracture energy calculation. X. Zhao
et al. [191] have employed a strain-to-fail test to measure the interface fracture toughness based
on buckle delamination mechanisms. Their method circumvents the complex modelling of
plastic deformation of the substrate. In addition, the in-plane compression load applied to the
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
109
top coat during the strain-to-fail test can inhibit the crack deflection into the YSZ layer. The
energy release rate associated with the buckling-driven delamination is estimated using the
buckle radius, coating stress and modulus and is taken as the interfacial toughness [191].
However, this requires that the buckle propagates stably without any ridge cracks. Failure to
fulfil these conditions will lead to an under-estimation to the interfacial toughness.
This chapter has two goals. Firstly, two single crystal superalloy blades (CMSX-4 and René
N5) coated with Pt-diffused γ/γ’ bond coats and YSZ top coats deposited by EB-PVD will be
investigated in terms of their cyclic oxidation behaviour in order to elucidate the effect of
superalloy substrate on TBC lifetimes. Secondly, the strain-to-fail test proposed by X. Zhao et
al. [191] is combined with the 3D-DIC technique to compare the coatings’ interfacial toughness
and its evolution with oxidation. The 3D-DIC technique can analyse the whole stable buckling
propagation process before any ridge crack initiation, thus allowing a reliable determination of
interface toughness values based on the well-established buckle-driven delamination
mechanism [139].
4.2 Experimental procedures
4.2.1 Sample preparation
The TBC blades investigated in this study were provided by Rolls-Royce plc. The composition
of the two superalloy substrates (CMSX-4 and René N5, denoted as X4 and N5, respectively)
determined by energy-dispersive X-ray fluorescence (EDXRF; PANalytical MiniPal 4) is given
in Table 4.1. These two superalloy blades were coated with Pt-diffused γ/γ’ bond coats
followed by the YSZ top coat. Firstly, the blades were grit blasted (alumina particles) and then
electroplated with Pt followed by annealing in vacuum at high temperatures
for Pt diffusion. The annealing resulted in a Pt-diffused bond coat ~ 30 µm thick. A ~ 175 µm
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
110
YSZ top coat (~ 7 - 8 wt.% Y2O3) was deposited on one side of the blades by EB-PVD using a
commercial process. A total of five blades for each type of superalloy substrate were fabricated.
Table 4-1 Superalloy compositions (atomic %) by EDXRF
Superalloy Ni Cr Al Ta Ti W Re Mo Co
René N5 Bal. 7.8 13.1 2.2 - 1.6 0.9 0.9 8.0
CMSX-4 Bal. 6.8 9.8 2.2 1.3 1.8 0.9 0.5 10.2
4.2.2 Thermal treatment
Cyclic oxidation was performed on some blades in laboratory air between room temperature
and 1200 °C in a cycle furnace (CM™). Each cycle consisted of 10 min ramping period, 1 h
holding time at 1200 °C and 10 min fan-assisted air quenching.
Other blades were exposed to isothermal oxidation at 1150 ºC in laboratory air in a box furnace
(CM™) for different periods of time up to 45 h. These isothermally oxidized specimens were
used for the strain-to-fail test for an evaluation of interface toughness. Isothermal oxidation at
a lower temperature (1150 ºC) can ensure that no sub-critical cracks initiation at the interface
after oxidation, thus the intrinsic interface toughness can be measured and compared for the
two TBCs with different substrates. Flat parts of the blades were cut by SiC blade into 6×6×1.5
mm samples using precision machine (Accutom 5, Struers) and used for the strain-to-fail test.
4.2.3 Microstructure characterization
For cross-sectional investigation of the samples, firstly they were ground by SiC paper up to
1200# and then polished using diamond paste to 0.25 μm finish. Scanning electron microscopy
(SEM; FEI Quanta 650) equipped with an energy dispersive X-ray spectrometer (EDS) was
used to examine the microstructure and composition of cross-sections. The bond coat surface
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
111
after spallation of the YSZ top coat was carbon coated and also examined by SEM.
Transmission electron microscope (TEM; FEI Talos F200A) with STEM (scanning
transmission electron microscope)/EDX (energy-dispersive X-ray) detector was used to
examine the TGO/coating interface. Focused ion beam (FIB; FEI Quanta 3D) in-situ lift-out
technique [168] was used to prepare cross-sectional site-specific specimens for TEM analysis .
4.2.4 Interface toughness measurement by the strain-to-fail test
4.2.4.1 Theoretical background
A strain-to-fail test combined with 3D-DIC technique was employed to measure the interface
toughness of EBPVD TBCs based on the buckle delamination mechanism. A sequence of
events involved in the buckling failure mode of a compressed coating includes [112]: ⅰ)
separations develop between the substrate and the coating; ⅱ) the film buckles once the
separation reaches a critical size; ⅲ) the buckle propagates along the coating/substrate interface;
iv) the crack deflects toward the coating surface leading to the coating spallation. Buckling of
a coating subject to an equi-biaxial compression, σ, occurs at a critical stress [192]:
σc=1.22[𝐸𝑐
1−𝜐𝑐2](ℎ/𝑏)
2 (4.1)
where b is the separation radius, h is the coating thickness, Ec and υc are the elastic modulus
and Poisson ratio of the coating, respectively. When the stress in the coating exceeds the critical
stress σc, the coating buckles away from the substrate and the energy release rate G is described
as [192]:
G=[1 + 0.9(1 − 𝜐𝑐)]−1𝐺0 [1 − (𝜎𝑐/𝜎)
2] (4.2)
where G0 is the elastic energy stored in the unbuckled coating per unit area, given by [192]
G0=1−𝜐𝑐
𝐸𝑐𝜎2ℎ (4.3)
The mode-dependent interface toughness Γ(ψ) is defined as [139]:
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
112
Γ(ψ)= Γicf(ψ) (4.4)
where Γic is the mode Ⅰ toughness and the criterion for a buckle to propagate is [112]
G= Γ(ψ) (4.5)
and
f(ψ)=1+ 𝑡𝑎𝑛2[(1 − 𝜆) ψ] (4.6)
where λ is a coefficient reflecting the interface friction (λ is about 0.3 for the bond coat/TGO
interface [112]). ψ is the loading phase angle, which represents the proportion of mode Ⅱ to
mode Ⅰ fracture, given by [139]:
tan ψ = (cos𝜔 + 𝜂sin 𝜔)/(ηcos𝜔 − sin𝜔) (4.7)
where 𝜔 is the phase angle shift generated by the elastic mismatch. For an alumina/Ni interface,
𝜔 = 52° [112].
And η is given by:
η = 0.25(1+𝜐𝑐)[(σ/𝜎𝑐 − 1)/g2]1/2 (4.8)
with
g2 = 0.25(1+𝜐𝑐) + 0.22(1−𝜐𝑐2) (4.9)
Therefore, if the stress σ of the coating, the buckling radius b and the elastic modulus Ec of the
coating are known, we can calculate the interfacial toughness Γ ic for the bond coat/TGO
interface based on the above expressions.
4.2.4.2 Strain-to-fail test coupled with 3D-DIC
Although both the TBC top coat and the alumina layer are in compression after thermal
exposure mainly due to the thermal mismatch between the ceramic layers and the metal
substrate, this stress is not sufficient to make the coating buckle from the substrate
spontaneously [191]. So an external compressive strain is applied to the pre-stressed coating to
facilitate the buckling process as schematically shown in Fig. 4.1. The compression load was
applied along the direction parallel to the TBC/substrate interface (y-axis direction in Fig. 4.1)
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by a Static Testing Machine Instron 5569H1549 at a loading rate of 0.07 mm∙min-1. The
buckling process of the EBPVD TBCs was monitored by two high speed cameras of the 3D-
DIC system (LaVision 3D DIC VP18-0021), at a frequency of 5 Hz (~200 ms exposure time
for each frame).
Fig. 4.1 Schematic view of the experimental setup: both high speed cameras take images of the
TBC coating surface (x-y plane) during the compression test, and the compression load is along
the y-direction.
4.2.4.3 Determination of the coating stress
The stress in the coating upon buckling is originally composed of both the residual stress and
the externally-applied stress. The applied stress can be calculated by the strain determined by
DIC, which will be illustrated later. The residual stress of alumina scale was measured by
luminescence spectrum on Renishaw Invia Raman system (RenishawTM, Gloucestershire, UK)
with an argon laser source (λ=633 nm) [193]. All measurements were performed through the
top coat surfaces. Before each experiment, the spectrometer was calibrated by taking a
spectrum from a standard pure silicon sample. The fluorescence spectra from Cr3+ in alumina
scale were collected for each measurement with five seconds acquisition time. Cr3+
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fluorescence in α-Al2O3 produces the R1 and R2 doublet at ~14400 cm-1. To determine the
peak positions in each spectrum, all spectra were deconvoluted in Wire 4.2 software
(RenishawTM) with an automatic fitting program by two mixed Gaussian-Lorentzian functions
[166]. The stress was estimated from the peak shift ∆𝜈 of the R2 line with respect to the
unstrained single crystal sapphire by assuming an equi-biaxial stress state and randomly
distributed TGO grains. The stress σ in the alumina scale was calculated by [130]:
∆𝜈 = 5.07(𝑐𝑚−1𝐺𝑃𝑎−1)σ (4.10)
4.2.4.4 Determination of the YSZ modulus
Due to the columnar structure of the EB-PVD TBCs, the stiffness of the coating can be highly
anisotropic. Since the energy release rate associated with the buckling-driven delamination
depends mainly on the coating’s in-plane elastic modulus, the nano-indentation test has been
used on the polished cross-sections of the TBCs to determine their in-plane elastic modulus. A
MTS XP nano-indenter was used. For each sample, a 2×10 (60 μm step size) array of indents
were made on the cross section using a Berkovich indenter with a penetration depth of 2 μm.
The Young’s modulus of YSZ top coat was calculated by the Oliver-Pharr method [194].
4.2.4.5 Measuring buckling radius by 3D-DIC
The recorded digital images were analysed using Davis 10.0 software. The displacement of the
coating surface was first determined by registering each frame with the reference image taken
prior to loading using a lease square matching (LSM) algorithm. In order to highlight the buckle,
the images showing the deformed coating surface were first corrected by the calculated
displacement field g (x + u (x)) and subtracted from the reference image f (x). The obtained
difference image (residual field) can clearly reveal cracks or any geometric features on the
sample during the loading process, which cannot be distinguished on as-deformed images [195].
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4.3 Results
4.3.1 Microstructure of the as-received TBCs
Fig. 4.2 As-received TBCs with different substrates, a) - d): cross-sectional SEM images; e)
and f): Ni, Al and Pt concentration profiles by SEM/EDX linescan along the lines in c) and d).
c) N5, as-received
e) f)
20 μm
a) N5, as-received b) X4, as-received
d) X4, as-received
100 μm
YSZ ~ 180 µm
BC ~ 27 µm
γ’ γ
EDX linescan
Alumina particles (grit blasting)
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Cross-sectional SEM images illustrating the as-received microstructure of the TBCs on two
different superalloy substrates are shown in Fig. 4.2. From Fig. 4.2 a and b, it can be seen that
both the YSZ top coat and the Pt diffusion bond coat (BC) were prepared to identical thickness
for the two superalloy substrates. The top coat had the characteristic columnar grain structure
typically produced by EB-PVD. As shown in Fig. 4.2 c and d, the Pt-diffusion bond coat
consisted of γ phase (darker contrast) and γ’ phase (brighter contrast). And the bond coat on
the X4 and N5 substrates both contained a grit-line (residual Al2O3 particles from the grit-
blasting process), which was at ~ 10 µm distance from the bond coat/top coat interface. The
EDX concentration line profiles of Ni, Pt and Al of the bond coats on two different substrates
are shown in Fig. 4.2 e and f, respectively. The Ni and Al concentration profiles on different
substrates were very similar, while the average Pt concentration of the bond coat on N5
substrate was slightly higher than that on X4 substrate.
4.3.2 Cyclic oxidation testing
4.3.2.1 YSZ lifetime
The lifetimes of TBCs with N5 and X4 substrates for 1-h cyclic test at 1200 ºC are illustrated
in Fig. 4.3. For each substrate, a total of three blades were cycled. The failure of the coating
was defined as ~ 20% spallation of the top coat (only consider the flat part of the blade). It can
be seen that the lifetime deviation for the three blades of the same substrate was small.
Moreover, three blades of the X4 substrate all exhibited longer lifetimes (30 cycles in average)
compared to that of the N5 substrate (24 cycles in average). The average lifetime of YSZ
coating on the X4 substrates reported here were similar to other studies on the same TBC
system under the same cycling condition [182, 196]. However, the higher lifetimes of TBCs
on X4 substrates than N5 substrates were in conflict with the result by B. A. Pint et al. [186],
which concluded that the Pt diffusion bond coat on superalloy X4 exhibited the shortest YSZ
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coating lifetime among different superalloy substrates including N5. Although the reason for
this conflict result is still not clear, it might be related to the differences in the coating
preparation process.
Fig. 4.3 YSZ lifetime for the cyclic oxidation testing of TBC-coated different superalloy blades.
4.3.2.2 Microstructural evolution
Fig. 4.4 exhibits the cross-sectional SEM images and concentration profiles of Ni, Pt and Al in
the bond coat on N5 and X4 substrates, after 5 1-h cycles at 1200 °C. At this stage, cracks along
the TGO/bond coat interface were observed on the N5 substrate (Fig. 4.4 a), whereas no
interfacial crack was found for the coating on the X4 substrate (Fig. 4.4 b). This indicates that
the loss of adhesion between the bond coat and TGO was earlier for TBC on the N5 substrate
than X4 substrate, which coincides with the longer lifetime of the TBC on the X4 substrate
(Fig. 4.3). As can be seen from Fig. 4.4 c and d, a γ-phase zone arose near the bond coat surface
for both TBC systems as a result of Al depletion. This is ascribed to the interdiffusion between
the substrate and the bond coat and aluminium loss due to oxidation. It is notable that the Al
concentration near the bond coat surface on the N5 substrate was nearly identical to that on X4
substrate (~ 10 at.%), suggesting that the Al depletion with thermal cycling was similar for the
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bond coats on the X4 and N5 substrate. While the Pt concentration near the bond coat surface
on the N5 substrate (~ 15 at.%, Fig. 4.4 e) was slightly higher than that on the X4 substrate (~
11 at.%, Fig. 4.4 f), as also found for the as-received coatings (Fig. 4.2 e and f).
Fig. 4.4 Cross-sectional SEM images: a) b) secondary electron (SE) mode; c) d) backscattered
electron (BSE) mode; e) f) corresponding SEM/EDX linescan elemental concentration profile
along the red lines in c) and d), respectively after 5 1-h cycles at 1200 °C.
N5, 5 cycles
c)
20 μm
e) f)
d)
b)
γ-phase, Al depletion zone
EDX linescan
50 μm
X4, 5 cycles
a)
Crack along the TGO/bond coat interface
No interfacial crack at this stage
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Fig. 4.5 a) b): cross-sectional SEM images (SE mode) and c) d): Ni, Al and Pt concentration
profiles by SEM/EDX linescan along the red lines in a) and b), respectively after 10 1-h cycles
at 1200 °C.
After further thermal cycling, as shown in Fig. 4.5, the cracks still propagated predominately
along the TGO/bond coat interface for the N5 substrate (Fig. 4.5 a). For the bond coat on the
X4 substrate, on the contrary, cracks along the TGO/bond coat interface (the red arrow in Fig.
4.5 b) and within TGO (black arrows in Fig. 4.5 b) co-existed after 10 cycles. Again, the
elemental concentration profiles for bond coats on the X4 and N5 substrates at this stage were
still similar when comparing Fig. 4.5 c and d, except for the slightly higher Pt concentration of
N5 substrate near the bond coat surface. Previous studies [175, 197] have pointed out that
higher Pt concentration near the bond coat surface can promote the uphill diffusion of
aluminium to the bond coat/TGO interface. However, it is observed here that the Al
concentration near the bond coat surface is quite similar for the coatings on N5 and X4
a)
c)
Cracks along TGO/BC interface
d)
50 μm N5, 10 cycles X4, 10 cycles
Cracks within TGO
b)
Cracks along TGO/BC interface
EDX linescan
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substrates. This suggests that Al depletion cannot explain the lifetime difference and there
should be some other reason for the shorter lifetime of the N5 superalloy system, which will
be discussed later.
Typical morphologies of the bond coat surface exposed by spallation of the top coat on the X4
and N5 substrate are shown in Fig. 4.6 a and b, and Fig. 4.6 c and d show the underside of a
spalled piece of the top coat. The backscattered electron (BSE) images can clearly demonstrate
the alumina (TGO), YSZ and exposed bond coat surface by different contrast. For TBC on the
N5 substrate, the exposed surface by failure of the coating showed mainly the bond coat (grey
area in Fig. 4.6 a), and correspondingly, the underside of the spalled top coat piece was covered
mainly by the alumina (dark area in Fig. 4.6 c), with only a small portion of YSZ presented
(bright area in Fig. 4.6 c). Therefore, the failure for the TBC on the N5 substrate was
predominantly along the bond coat/TGO interface, which agrees with the above cross-sectional
observations.
In contrast, TBC on the X4 substrate has failed in a different way. From Fig. 4.6 b, it can be
seen that the exposed surface for this system exhibited a mixture of alumina (dark area), YSZ
(mid-tone area, surrounded by alumina) and the metallic bond coat (bright area), suggesting
that the fracture for this system can take place at the TGO/bond coat interface and also within
the TGO. Sometimes, cracks can also extend laterally into the YSZ where TGO intrusions into
the bond coat existed (examples of this are shown in Fig. 4.5 b), thus leaving some YSZ
particles on the exposed bond coat surface. This is also consistent with the morphology of the
underside of spalled YSZ shown in Fig. 4.6 d.
In summary, the failure of TBC occurred predominantly along the bond coat/TGO interface for
TBCs on the René N5 substrate, whereas for TBCs on the CMSX-4 substrate, a mixed failure
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path (at the bond coat/TGO interface and within TGO) was observed. These observations are
self-consistent and also in agreement with the failures reported by other researchers [130, 185]
for similar specimens.
Fig. 4.6 a) and b): bond coat surface BSE images exposed by spalling after 28 cycles at 1200 °C;
c) and d): BSE images of the back side of spalled YSZ coating after 28 cycles.
Fig. 4.7 a shows the average TGO thickness evolution during the cycling test at 1200 °C. It can
be seen that the average TGO thicknesses on the X4 and N5 substrate were quite similar up to
20 cycles, because microstructural studies confirmed that the bond coat on both superalloy
substrates (X4 and N5) was able to develop a pure alumina layer without formation of other
oxides (e.g. Ni-oxides) during the cycling test until failure. In addition, although the average
lifetime of TBC on the N5 substrate was 24 cycles, in comparison with 30 cycles for the X4
a) N5, bond coat surface
b) X4, bond coat surface
d) X4, spallaed YSZ
underside c) N5, spallaed YSZ underside
Alumina
Bond coat
YSZ
Bond coat
50 μm
alumina
Alumina
YSZ
Alumina
YSZ
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substrate, both superalloy systems exhibited about the same TGO thickness near the time of
failure. This comparable TGO thickening rate indicates that the TBC lifetime difference of the
two superalloy substrates is not correlated to the TGO thickening evolution.
Fig. 4.7 b illustrates the root mean square roughness evolution of TGO/BC interface of the
TBCs on the X4 and N5 substrate, respectively. Clearly, similar interfacial roughness evolution
throughout the cycling test for different substrate systems indicates that the influence of TGO
rumpling on the observed lifetime difference can be neglected. B. A. Pint et al. [186] also
concluded that TGO rumpling had no effect on TBC lifetime by investigating the Pt-modified
aluminide coating on the X4 and N5 substrate, respectively.
Fig. 4.7 a) TGO thickness evolution during the cycling test; b) root mean square roughness
evolution of TGO/BC interface by processing of digitized profiles.
Overall, the cycling test showed that the TBC on the CMSX-4 superalloy substrate has a 25%
longer lifetime compared to that on the René N5 substrate. No significant difference regarding
interfacial morphology and TGO growth rate evolution has been observed for the TBCs on
different substrates, which rules out these effects on the observed lifetime difference. The
microstructural investigations showed that interfacial cracking occurred earlier for the TBC on
the N5 substrate than that on X4 substrate. In addition, the fracture was mainly found along the
a) b)
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bond coat/TGO interface for the René N5 specimens, whereas for the TBC on the CMSX-4
substrate, a mixed failure path (at the bond coat/TGO interface and within TGO) was observed.
This strongly indicated that the bond coat/TGO interfacial degradation was faster and to a
greater extent for the N5 specimen, as compared to the X4 specimens under identical cyclic
exposure. To confirm this, a strain-to-fail test based on a previously established model was
employed to measure the bond coat/TGO interface fracture toughness evolution during
isothermal oxidation of TBCs with the N5 and X4 substrates, respectively. The results are
presented as follows.
4.3.3 Strain-to-fail compression test coupled with 3D-DIC
DIC residual fields were utilised to reveal the whole buckling propagation process during the
strain-to-fail test. Fig. 4.8 exhibits a) the reference image f (x), b) the deformed images
corrected by the DIC displacement field g (x + u (x)) with increasing loading time and c) the
corresponding residual field Φ (x) (differences between g (x + u (x)) and f (x)). Buckling can
hardly be seen on the deformed images (Fig. 4.8 b), but appeared clearly on the residual fields
(Fig. 4.8 c) due to the luminosity change induced by the buckling initiation. Buckling initiated
at the edge of the analysing area (not the sample edge, as indicated by the optical image of the
sample surface in Fig. 4.8 a), as shown by a red ellipse in Fig. 4.8 c, and then it grew stably.
Based on this image processing method, the buckling radius and the corresponding strain can
be correlated during the stable buckling propagation process. This provides a more accurate
buckling measurement compared to the previous work by X. Zhao et al. [191] as their work
only considered the final buckled stage.
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Fig. 4.8 a) Optical image taken by the camera for DIC showing the sample surface prior to
applying the load. The region-of-interest (ROI) is highlighted with the red rectangle and is used
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as the reference image f (x); b) deformed images at several loading scales corrected by a DIC
displacement field g (x + u (x)) and c) the corresponding residual field Φ (x) of a sample (as-
received TBC with N5 substrate) during the test. The red ellipse in c) highlights the occurrence
of buckling.
Buckles at intermediate propagation stage (e.g. the buckle shown by a red rectangle in Fig. 4.8
c) were considered for the calculation of interface toughness of all tested specimens. Fig. 4.9 a
shows the buckling radii of TBCs with N5 and X4 substrates, respectively, as a function of
oxidation time. It can be seen that the buckling radii for TBCs with two different substrates
were similar up to 45 h oxidation time. Fig. 4.9 b) shows the strains for corresponding buckling
radii calculated by DIC of TBCs with N5 and X4 substrates, respectively. The compressive
strains of both TBCs have decreased in magnitude with increasing oxidation time. This might
indicate degradation of coating adhesion during oxidation. These buckling strains and radii will
be used for the calculation of interface toughness.
Fig. 4.9 Evolution of a) average buckling radii, and b) corresponding strains calculated by DIC
as a function of oxidation time at 1150 °C for specimens with the X4 and N5 substrate,
respectively.
The compressive residual stresses in the TGO of isothermal oxidized specimens with N5 and
X4 substrates are shown in Fig. 4.10 a). The TGO stress on N5 substrate was ~ -2.2 GPa for
a) b)
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as-received samples and reached a peak of ~ -2.8 GPa after 30 h oxidation, then the stress
slightly decreased to ~ -2.6 GPa after 45 h oxidation. This stress trend is the same as a previous
investigation on TBCs with Pt-diffused γ/γ’ bond coats under cyclic oxidation [198]. The TGO
stresses for the specimens with X4 substrates exhibited a similar trend but the stress values
were smaller than that of N5 substrates. The average TGO thicknesses are plotted in Fig. 4.10
b) as a function of oxidation time for two TBC systems. The TGO thicknesses for as-received
TBCs with N5 and X4 substrates were ~ 0.5 μm, and then the TGO growth followed a parabolic
law and grew to ~ 3.0 μm after 45 h.
Fig. 4.10 a) Residual stress of TGO and b) average TGO thickness as a function of isothermal
oxidation time at 1150 °C for TBCs with the X4 and N5 substrate, respectively.
The elastic moduli of YSZ top coat measured by the cross-sectional nano-indentation are listed
in Table 4.2. The elastic moduli increased slightly with increasing oxidation time. This is due
to the sintering of YSZ during the heat treatment.
Table 4-2 Elastic moduli of YSZ top coat measured by the cross-sectional nano-indentation
test
Oxidation time (h) 0 15 30 45
EYSZ (GPa) 25.6 + 1.7 27.1 + 2.0 30.1 + 2.3 32.8 + 3.1
a) b)
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4.4 Discussion
4.4.1 Estimation of the interfacial toughness for TBCs on N5 and X4 substrates
So far, all parameters required for estimation of interfacial toughness are presented. Then the
interfacial toughness for TBCs on the N5 and X4 substrates can be estimated based on the
buckling scenario as described in Section 4.2.4.1. However, this buckling principle is for a
single layer coating system, while the coating here is actually a bilayer consists of a TGO layer
and a YSZ top coat. Therefore, some parameters will be modified to make the calculation
suitable for this bilayer. The (TGO + YSZ) bilayer can be approximated as a single layer, with
the effective Young’s modulus given by [146]:
Eeff = 𝐸𝑌𝑆𝑍ℎ𝑌𝑆𝑍/(1−𝜐𝑌𝑆𝑍)+𝐸𝑇𝐺𝑂ℎ𝑇𝐺𝑂/(1−𝜐𝑇𝐺𝑂)
ℎ𝑌𝑆𝑍+ℎ𝑇𝐺𝑂 (4.11)
where
𝐸𝑌𝑆𝑍, 𝐸𝑇𝐺𝑂 -Young’s modulus of the YSZ layer and the TGO layer,
ℎ𝑌𝑆𝑍, ℎ𝑇𝐺𝑂 - Average thickness of the YSZ and TGO layer, and
𝜐𝑌𝑆𝑍, 𝜐𝑇𝐺𝑂 - Poisson’s ratio of the YSZ and TGO layer, respectively.
The effective stress of this bilayer can be written as [146]:
σ = 𝜎𝑌𝑆𝑍ℎ𝑌𝑆𝑍+𝜎𝑇𝐺𝑂ℎ𝑇𝐺𝑂
ℎ𝑌𝑆𝑍+ℎ𝑇𝐺𝑂 (4.12)
where σ is the stress of each layer, given by [191]:
σYSZ = σYSZ,0 + σYSZ, applied = 𝐸𝑌𝑆𝑍∆𝛼∆𝑇
1−𝜐𝑌𝑆𝑍 +
𝐸𝑌𝑆𝑍𝜀
1−𝜐𝑌𝑆𝑍 (4.13a)
σTGO = σTGO,0 + σTGO, applied = σTGO,0 +𝐸𝑇𝐺𝑂𝜀
1−𝜐𝑇𝐺𝑂 (4.13b)
where the subscripts ‘0’ and ‘applied’ refer to the stress before and after applying the
compressive strain. ε is the buckling strain, as illustrated in Fig. 4.9 b. ∆𝛼 is the difference of
coefficient of thermal expansion between the YSZ and the substrate (~ 4×10-6 K-1 [4]); ∆𝑇 is
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128
the temperature drop from the oxidation temperature (1150 °C) to the room temperature (25 °C).
σTGO,0 is the residual stress of the TGO scale, measured by luminescence spectrum (Fig. 4.10
a), and the residual stress of the YSZ layer is calculated by assuming the stress to be thermally
elastic, i.e. ε = 0 in Eq. 4.13a.
For the calculation of Γic (mode Ⅰ interfacial toughness), the buckling strains and corresponding
radii were taken from Fig. 4.9, and the stress of the TGO and YSZ bilayer was calculated
according to Eqs. 4.12 and 4.13. The TGO stresses and average thicknesses were taken from
Fig. 4.10. The other parameters were 𝐸𝑇𝐺𝑂 = 380 GPa, υTGO = υYSZ =0.22, hYSZ= 180 µm and
elastic moduli of YSZ coating shown in Table 4.2.
The calculated interfacial toughness for TBCs on the N5 (black squares) and X4 (red circles)
substrates as a function of oxidation time is shown in Fig. 4.11. The mode Ⅰ interfacial
toughness values of TBCs reported by other researchers [132, 146, 150] are also indicated,
which were comparable to the Γic values in this study when considering the phase angle effect
in different test methods. There are two noteworthy characteristics of the toughness values of
the two different TBCs. Firstly, for both TBC systems, the interfacial toughness decreased with
increasing oxidation, indicating that the bond coat/TGO interface was weakened by oxidation.
This behaviour agrees with some previous studies [146, 188]. However, it is noticed that a
recent study [132] has found that no decrease of interfacial fracture toughness with thermal
cycles up to 100 cycles at 1150 °C. The cause of this difference between different test methods
is unclear at this stage and further work is needed to clarify this. Another feature is that although
the mode I interfacial toughness of as-received TBCs were very similar for the two substrates,
the interfacial toughness of N5 samples decreased much faster with the oxidation time than that
of the X4 samples. For instance, the mode Ⅰ toughness of TGO/bond coat interface for the as-
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
129
received N5 specimen was ~ 33 J/m2 and it dramatically decreased to ~ 12 J/m2 after 15 h
oxidation. While for TBCs with X4 substrates, the interfacial toughness for as-received
samples was ~ 30 J/m2 and reduced to ~ 21 J/m2 after 15 h. The reason for this will be given in
Section 4.4.2.
Fig. 4.11 Mode Ⅰ interface toughness of TBCs on N5 (black square) and X4 substrates (red
circle) as a function of oxidation time, respectively. Data of other TBC systems are also
included for comparison: Pt diffusion bond coat (1150 °C, X. Wang et al. [132]); NiCoCrAlY
bond coat (Yu-Fu Liu et al. [150]); β-NiPtAl bond coat (Vasinonta and Beuth [146]). Note that
the data in all references has been replotted by considering the different phase angles in
different test methods.
4.4.2 Interface degradation of TBCs on different substrates
In order to find out the reason for the faster degradation of TGO/bond coat interface on the N5
substrate, high resolution STEM/EDX mapping analysis was carried out. The results for TBCs
with the N5 and X4 substrate after 3 1-h cycles at 1200 °C are shown in Fig. 4.12 and 4.13,
respectively. As can be seen from Fig. 4.12 a, pre-failure cracking was observed for TBC on
CHAPTER 4 EFFECT OF SUPERALLOY SUBSTRATE ON THE LIFETIME OF EBPVD TBCs
130
the N5 substrate. Moreover, sulfur was found to segregate at the bond coat/TGO interface, as
shown by Fig. 4.12 c. This sulphur segregation has been widely reported to be detrimental to
the coating performance because sulphur can significantly degrade the interfacial toughness
[74, 199]. In contrast, no cracks existed on the specimen for X4 substrate after 3 cycles (Fig.
4.13 a), and no sulfur segregation was observed at the TGO/bond coat interface (Fig. 4.13 c).
The STEM/EDX results can explain the faster degradation of TGO/bond coat interface
toughness on the N5 substrate, compared to that of X4 substrate [200]. Further research is
required to find out the reason for this different sulphur segregation behaviour for these two
superalloy substrates.
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Fig. 4.12 High resolution STEM/EDX analysis of the TGO/bond coat interface: a) and b)
HAADF (high angle angular dark field)/STEM image, b) is the red box area shown in a); c) -
g) STEM/EDX mapping of TGO on N5 substrate after 3 cycles at 1200 °C.
Fig. 4.13 High resolution STEM/EDX of the TGO/bond coat interface: a) and b)
HAADF/STEM image, b) is the red box area shown in a); c) - g) STEM/EDX mapping of TGO
on X4 substrate after 3 cycles at 1200 °C.
H. M. Tawancy et al. [180] pointed out that for TBCs on Ti-containing substrates, the
interfacial adhesion can be degraded by the segregation of TiO2 particles at the TGO/bond coat
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132
interface, resulting in poor TBC performance. However, in this study no Ti-oxide was found
within the TGO or at the bond coat/TGO interface for the Ti-containing CMSX-4 substrates.
This indicates that Ti in the superalloy substrate is not necessarily detrimental to the
performance of Pt-containing bond coats, as also suggested by other researchers [167, 201]. It
has been proposed that Pt in the γ/γ’ bond coat can serve as a sink for Ti [202], which might
explain this annihilation of Ti precipitates observed here.
4.5 Summary
In this contribution, cyclic oxidation tests at 1200 °C have been carried out on TBCs with the
CMSX-4 and the René N5 substrate, respectively. And a strain-to-fail method combined with
the 3D-DIC technique was employed to measure the bond coat/TGO interface toughness of
TBCs with different substrates. Main findings can be summarized as follows:
• TBCs based on the CMSX-4 superalloy have a 25% longer average lifetime compared
to that on the René N5 superalloy during the cyclic test. No significant difference
regarding interfacial morphology and TGO growth rate has been observed for these
TBCs with different substrates until failure of the coating.
• Spallation of the TBC occurred mainly along the bond coat/TGO interface for TBC
with the René N5 substrate, whereas for TBC with the CMSX-4 substrate, a mixed
failure path (along the bond coat/TGO interface and within TGO) is observed.
• The strain-to-fail test showed that although the mode I interfacial toughness (Γic) values
were almost identical for the two substrates in the as-deposited state, the interfacial
toughness of René N5 specimens decreased faster with increasing oxidation time.
• This faster degradation of TGO/bond coat interfacial toughness for the N5 substrate can
be ascribed to the sulfur segregation at this interface, which has been confirmed by high
resolution STEM/EDX.
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Chapter 5 The Al-enriched γ’-Ni3Al-base bond coat for thermal
barrier coating applications
5.1 Introduction
The quest for increased operating temperatures of gas-turbine engines, in order to achieve a
higher engine efficiency, has driven the development of thermal barrier coatings (TBCs) for
protections of superalloy components in the hot section of gas-turbine engines [51, 73]. A TBC
system consists of a thermally-insulating ceramic top coat, an intermediate metallic bond coat
and the underlying superalloy substrate. The bond coat has been considered as the most crucial
part of a TBC system which has two primary functions. Firstly, it can provide adhesion between
the ceramic top coat and the metallic substrate. Maintaining the adhesion of top coats to the
superalloy substrates is vital for the lifetime of superalloys in the high temperature combustion
environment [203]. Secondly, the bond coat serves as an aluminium reservoir from which the
coating can form and maintain a slow-growing α-Al2O3 thermally grown oxide (TGO) layer.
This TGO layer can protect the underlying superalloy substrates from being oxidized at high
temperatures.
The diffusion bond coats based on the Pt-modified β-NiAl intermetallic compound, have been
optimized to form and maintain a dense α-Al2O3 layer. This ensures the coating to provide
good oxidation resistance for the underlying superalloys. However, the Pt-modified β-NiAl
coating (denoted as the β-NiPtAl coating hereinafter) exhibits inferior mechanical properties,
especially at high temperatures. For example, its strength decreases sharply at high
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temperatures, thus makes it vulnerable to the deformation (referred to as rumpling) during
cyclic oxidation [204]. Rumpling of the bond coat surface can cause the detachment of the top
coats from the bond coat, leading to the subsequent spallation. Because the ceramic top coat
has low out-of-plane compliance which prevents it from deforming together with the bond coat
[203].
Another type of diffusion bond coat, the Pt-diffused γ-Ni/γ’-Ni3Al coating, has attracted
increasing attention. Early work by Tatlock et al. [205] confirmed that Pt can improve the
oxidation resistance of alloys with γ+γ’ compositions. Recent progress [175, 206] showed that
the Pt-diffused γ/γ’ coating exhibits a significant reduction of rumpling after high temperature
cyclic exposures, compared to the β-NiPtAl coating. This benefit can be ascribed to the high
strength of the γ’- Ni3Al structure at elevated temperatures [207, 208]. Despite these advantages,
the low Al content (~ 20 at. %) of the Pt-diffused γ/γ’ coatings has raised concerns regarding
their long-term oxidation behaviour [209]. Several studies [30, 52, 210] have reported the
spinel (NiAl2O4) formation in the TGO of this coating during prolonged oxidation because of
the insufficient Al source. The brittle spinel can significantly compromise the interface
adhesion between the TGO and the top coat, leading to spallation of the top coat.
Overall, there are both advantages and disadvantages for the two types of Pt diffusion coatings.
Although the β-NiPtAl coating can provide effective oxidation protection, it degrades
mechanically by rumpling during cyclic oxidation. The Pt-diffused γ/γ’ coatings exhibit
excellent rumpling resistance during cyclic exposures but has inferior oxidation resistance due
to low Al contents. Another critical issue with these two diffusion coatings is the coating-
substrate interdiffusion, which can also result in coating failure [74]. Pt addition was initially
intended to mitigate this interdiffusion and promote the outward Al diffusion to the
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coating/TGO interface [43]. However, Chen and Little [211] have studied the degradation of
β-NiPtAl coatings on the CMSX-4 substrate and found that Pt neither inhibited coating-
substrate interdiffusion nor promoted the outward Al diffusion. Other studies have also
reported that Al depletion due to interdiffusion can result in the coating degradation for the Pt-
diffused γ/γ’ coatings [30].
Therefore, the objective of the present study is to fabricate a Pt diffusion coating that can resist
high-temperature rumpling while maintaining adequate oxidation resistance and exhibiting
mitigated coating-substrate interdiffusion. The proposed Al-enriched pure γ’-phase coating
(denoted as the Al-enriched γ’-phase coating) not only maintains the rumpling resistance of the
γ’-Ni3Al structure, it also shows comparable oxidation performance to the β-NiPtAl coating.
In addition, this new γ’-phase coating exhibits much less pronounced coating-substrate
interdiffusion during oxidation compared to the γ/γ’ two-phase coating, which contributes to
its superior TGO spallation resistance. There are two steps for fabricating the Al-enriched pure
γ’-phase coating. Firstly, a Pt-diffused intermediate coating will be fabricated on the CMSX-4
superalloy by a selective γ-etching process and subsequent Pt electroplating [24]. Then a pack
cementation aluminizing process will be carried out on this Pt-diffused intermediate coating to
obtain the Al-enriched pure γ’-phase coating. In addition, industry-standard Pt-diffused γ/γ’
coatings and β-NiPtAl coatings will be investigated and compared to this new Al-enriched γ’-
phase coating in terms of both isothermal and cyclic rumpling behaviours. The focus of this
manuscript will be on the coating-substrate interdiffusion, TGO microstructure & spallation
and the rumpling behaviour of these coatings.
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5.2 Experimental procedures
5.2.1 Sample preparation
The Al-enriched γ’-phase coatings were fabricated at the University of Manchester. Standard
Pt-diffused γ/γ’ coatings and β-NiPtAl coatings were supplied by Rolls-Royce®. The next
section will describe the fabrication of Al-enriched γ’-phase coatings in details, followed by a
brief introduction of the fabrication of the industry-standard coatings.
5.2.1.1 Fabrication of Pt-diffused intermediate coatings
The Pt-diffused intermediate coatings were applied to CMSX-4 single crystal Ni-based
superalloy (Table 5.1, Rolls-Royce plc) substrates. Cylindrical CMSX-4 bars (20 mm in
diameter) with the <001> orientation aligned with the cylinders’ long axis were cut into buttons
of 5 mm thickness by a SiC blade using a precision machine (Accutom 5, Struers). All buttons
were ground using 400# SiC paper, and then cleaned by acetone in an ultrasonic bath for 20
min. Subsequently, the buttons were selectively electrolyte etched to remove the γ matrix of
the single crystal superalloy. The selective electrolyte γ-etching parameters are shown in Table
5.2 [24]. Fig. 5.1 exhibits a) the as-etched cross section by the focus ion beam (FIB; FEI Quanta
3D) system coupled with SEM, b) the as-etched surface morphology and c) etching time -
thickness plot, respectively. It can be seen from a) and b) that γ phases have been totally
removed and the cuboidal γ’ precipitates were retained after the selective γ-etching. Fig. 5.1 c)
shows a linear relationship of etching time vs. thickness (~ 5.0 µm/min).
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Table 5-1 Composition of CMSX-4 superalloy
Element Ni Al Cr Co Ta Ti W Re
Wt. % 61.4 5.6 6.4 9.6 6.6 1.0 6.4 2.9
Fig. 5.1 a) cross-sectional FIB/SEM image and b) surface image of as-etched samples; c)
etching time-thickness plot and d) cross-sectional FIB/SEM image after Pt electroplating on
the etched substrate.
The as-etched substrates were Pt-electroplated where Pt filled the gaps left by removing the γ
matrix and then formed an overlay layer above the sample surface (Fig. 5.1 d). After Pt
electroplating, all samples were immersed in hot distilled water (80°C, 1h) to remove the
remaining salts (from the electroplating bath) on the sample surface, followed by vacuum heat
treatment at 1150°C for 2 h to form the Pt-diffused intermediate coatings.
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Table 5-2 Electrolyte γ-etching bath parameters
Chemical formula 10 vol.% H3PO4 + 90 vol.% H2O
pH 0.04-0.06
Temperature Room temperature
Cathode current density 120 mA/cm2
Magnetic stirring speed 120/min
5.2.1.2 Fabrication of Al-enriched γ’-phase coatings by pack cementation
A low-temperature pack cementation aluminizing process was carried out on the Pt-diffused
intermediate coatings to increase the Al content. Firstly, the powder mixture of aluminium (The
Aluminium Powder Co. Ltd), CrCl3 activator (Sigma-Aldrich) and Al2O3 inert filler powder
(Honeywell FlukaTM) were manually ground with a mortar and a pestle, followed by
mechanical mixing of the powder mixture for 4 h. Then the Pt-diffused intermediate coating
samples were buried in a crucible with the powder mixture and sealed by the cement. The
following heat treatment was conducted in a tube furnace with Ar gas flowing at temperatures
650 to 950 °C to aluminize. Finally, samples were annealed in vacuum at 1150°C for 4 h to
form the Al-enriched γ’-phase coatings. Different pack cementation parameters (powder
mixture composition, heat treatment temperature & time, etc.) were investigated in order to
obtain Al-enriched pure γ’-phase coatings, which will be detailed in section 5.3.1. The
fabrication process of the Al-enriched γ’-phase coating is summarized in Fig. 5.2. It is noted
that the average Al concentration of the Pt-diffused intermediate coating (~ 10 - 12 at. %) is
lower than that of the Pt-diffused γ/γ’ coatings (~ 16 - 19 at. %) due to the etching process
which removes all Al contents of the γ phase, thus ‘diluting’ the Al content of the coating.
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Fig. 5.2 Two steps to fabricate the Al-enriched pure γ’-phase coating: Ⅰ. Fabricate Pt-diffused
intermediate coatings; Ⅱ. Pack cementation aluminizing on the intermediate coating.
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The two industry-standard coatings were also applied to the CMSX-4 substrates. The Pt-
diffused γ/γ’ coatings were produced on grit blasted substrates by Pt electroplating followed
by a diffusion anneal in vacuum [196], and β-NiPtAl coatings were produced by Pt
electroplating followed by a low-activity chemical vapour deposition (CVD) aluminizing [35].
5.2.2 Thermal treatment
Isothermal oxidation was performed at 1150ºC in laboratory air in a box furnace (CM™). The
samples were oxidized for different periods of time up to 100 h, followed by air quenching.
Cyclic oxidation was performed in laboratory air between room temperature and 1150°C in a
cycle furnace (CM™) in order to investigate the rumpling behaviour of the coatings. Each
cycle consisted of 10 min ramping period, 10 min holding time at 1150°C and 10 min fan-
assisted air quenching.
5.2.3 Characterization methods
To identify the as-diffused coating phases, electron backscatter diffraction (EBSD,
NordlysNano, Oxford Instruments) was used in addition to X-ray diffraction (XRD, Philips
X'Pert) with Cu Kα radiation (λ=0.154 nm).
After oxidation, the optical microscope (Olympus BH2-UMA) was used to examine the surface
morphology of the oxides. XRD with Cu Kα radiation (λ=0.154 nm) was used to examine the
oxide phases. Then the site-specific cross-sectional images of the oxides were captured by SEM
coupled with FIB (FEI Quanta 3D). Finally, the cross-sections of the samples were ground by
SiC paper up to 1200# and then polished using diamond paste to 0.25 μm finish. SEM (FEI
Quanta 650) equipped with an energy dispersive X-ray spectrometer (EDS) was used to
examine the cross-sections for elemental diffusion profiles of the coatings.
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To investigate the rumpling behaviour of the coatings, firstly, Vickers micro-hardness
indentations (Duramin hardness tester, Struers) were placed into the surfaces of each coating
before the cyclic oxidation. These indentations can act as markers so that the rumpling
evolution of the same surface area on each sample could be tracked during the cyclic oxidation.
The coatings were removed from the cyclic furnace at specific cycles for surface roughness
characterization using an optical profilometer (Bruker). With the help of the indentation
markers, the digital images of the same area in the form of surface height, Z, as a function of
position x and y were recorded for each coating during the cyclic oxidation. The surface
rumpling magnitude was characterized by the root mean square roughness Rq:
Rq= √1
𝑛∑ (𝑍𝑖− 𝑍)
2𝑛𝑖=1 (5.1)
where Zi is the height of each point, 𝑍̅ is the average height of all points and n is the number of
total points.
5.3 Results
5.3.1 Synthesis of Al-enriched γ’-phase coatings by pack cementation
To find the appropriate pack cementation parameters for obtaining the Al-enriched γ’-coatings,
different pack cementation parameters have been tried with the guidance of Ni-Pt-Al phase
diagram to find the target concentration for the Al-enriched γ’-phase coating. Different pack
cementation parameters (powder composition, temperature and holding time) of 5 trial samples
(number 1-5) are given in Table 5.3. The coating compositions corresponding to sample 1-5
were measured by SEM/EDX and Table 5.3 lists the average Al, Pt and Ni concentrations of
the as-fabricated coating for each sample.
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Table 5-3 Different pack cementation parameters and the resulting coating composition by
EDS after vacuum anneal
Sample number 1 2 3 4 5
Powder
composition
(wt. %)*
CrCl3 (4.0)
+ Al (5.0)
CrCl3 (4.0) +
Al (5.0)
CrCl3 (1.0) +
Al (1.0)
CrCl3 (1.0) +
Al (1.0)
CrCl3 (1.0) +
Al (1.0)
Temperature
(°C)
950 800 800 650 650
Hold time at
temperature
(min)
<0.5 <0.5 <0.5 15 5
Average Al
concentration
(at.%) of γ’-
coating by EDS
35.9 23.3 31.0 33.5 19.2
Ni concentration
(at.%) of γ’-
coating by EDS
26.3 32.6 30.5 29.7 40.1
Pt concentration
(at.%) of γ’-
coating by EDS
29.8 29.5 30.6 31.2 29.2
*Al2O3 powder with balance weight.
Fig. 5.3 exhibits the Ni-Pt-Al phase diagram at 1150°C [202], and different dots represent the
coating compositions of sample 1-5. Compositions of the sample 1-4 are in the three-phase
region (Fig. 5.3) including the β-NiAl phase, which indicates that these samples were over-
aluminized. The inset SEM image in Fig. 5.3 shows the coating microstructure of the sample
3. It can be seen that the interdiffusion zone of this coating had some precipitates, which is
characteristic for a β-NiAl-base diffusion coating [108]. Sample 5 (aluminized below 700°C)
located at the γ-Ni+α-NiPt two-phase region, indicating this sample was under-aluminized. In
order to find the target concentrations for the Al-enriched γ’-phase coating, the Ni-Pt-Al phase
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diagram at the annealing temperature (1150°C) [202] was utilised, as shown in Fig. 5.4. Since
the Pt electroplating process was identical for all samples, the coatings exhibited similar Pt
concentration (~ 29.0 - 32.0 at. %), as presented by the two blue lines in Fig. 5.4. The region
defined by the two blue lines has intersected with the pure γ’-phase region, thus the upper and
lower limit of the target Al concentration for the Al-enriched pure γ’-phase coating can be
determined by the highest and lowest point of this intersection region, respectively. The upper
(~ 27 at. %) and lower (~ 22 at. %) limits of the target Al concentration for the Al-enriched
pure γ’-phase coatings were marked by the two horizontal red lines in Fig. 5.4. Recalling the
Al concentrations of different packs in Table 3, the upper and lower Al concentrations located
between that of sample 4 and sample 5. Therefore, a holding time between sample 4 and sample
5 can fabricate the Al-enriched pure γ’-phase coating.
Fig. 5.3 Ni-Pt-Al phase diagram at 1150°C [202]. The compositions of sample 1-5 are marked
by the different dots, respectively. The inset SEM image shows the as-fabricated cross-sectional
microstructure of sample 3.
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Fig. 5.4 Ni-Pt-Al phase diagram at 1150°C [202]. The two horizontal red lines represent the
upper and lower limit of Al concentration for the pure γ’-phase coating, respectively.
5.3.2 Microstructure of the as-received coatings
By 10 min holding time at 650 °C and the same pack powder composition as sample 4, the Al-
enriched pure γ’-phase coating was fabricated. The average Ni, Pt and Al concentration of this
coating is shown by the red triangle in Fig. 5.4, which is located in the single γ’-phase region.
Fig. 5.5 shows the XRD results of the three as-fabricated coatings. The γ phase is a disordered
face centre cubic (fcc) solid solution, and the γ’ phase is an ordered fcc crystal structure (L12
superlattice). The β-NiAl phase is an ordered body centred cubic (bcc) structure (B2
superlattice). According to Fig. 5.5, the XRD patterns for as-fabricated Al-enriched γ’-phase
coating and the Pt-diffused γ/γ’ coating were quite similar. Because the diffraction peaks of γ
and γ’ locate very close to each other except for the additional peaks of γ’. It is noted that there
was no β-phase in the Al-enriched γ’-phase coating, which confirms that the pack cementation
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process was not over-aluminized. In addition, the pattern of the β-NiPtAl coating only showed
β peaks, while the other two coatings also exhibited peaks of Pt.
Fig. 5.5 XRD patterns of as-fabricated Al-enriched γ’-phase coatings, Pt-diffused γ/γ’ and β-
NiPtAl coatings.
Fig. 5.6 a - c exhibit the cross-sectional SEM (backscattered electron, BSE) images of three as-
fabricated coatings. Only γ’-phase presented in the Al-enriched γ’-phase coating (Fig. 5.6 a),
whereas γ (dark contrast) and γ’ phase (bright contrast) were intermixed with each other in the
Pt-diffused γ/γ’ coating (Fig. 5.6 b). The β-NiPtAl coating exhibited an outer β-NiPtAl layer
and an inner interdiffusion zone (IDZ) (Fig. 5.6 c). A large number of white contrast
precipitates rich in refractory elements (Ta, Mo, W) can be seen, especially in the IDZ [185].
Moreover, numerous pores were observed in the Al-enriched γ’-phase coating, especially in
the area near the surface, while the other two coatings exhibited a dense microstructure. The
formation of these pores can be explained by the Kirkendall effect as a result of the
interdiffusion of Al, Ni and Pt during the fabrication process.
The EBSD phase mappings of each coating are shown in Figs. 5.6 d - i. Fig. 5.6 d confirmed
that no β or martensitic phases were observed in the Al-enriched γ’-phase coating. A previous
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attempt to aluminize the Pt-diffused γ/γ’ based coatings on nickel superalloys has failed due to
the over-aluminizing resulting in the formation of β or martensitic phase on the coating surface
[59]. However, this study suggests that the low-temperature pack cementation process is
capable of aluminizing the γ’-base coating without altering the γ’-phase microstructure. It is
also noted that the lower part of the Al-enriched γ’-phase coating had the same crystal
orientation as the superalloy substrate (Fig. 5.6 g), indicated limited Pt solid solution in this
lower part. The grains of both the Pt-diffused γ/γ’ coating (Fig. 5.6 h) and β-NiPtAl coating
(Fig. 5.6 i) showed random orientations. In addition, it was observed that the average grain size
of the Pt-diffused γ/γ’ coating and the Al-enriched γ’-phase coating was comparable, while the
grain size of the β-NiPtAl coating was much smaller compared to the other coatings, especially
the surface part (nano-grains). This can be attributed to the large amount of recrystallization
during the coating manufacturing process of the β-NiPtAl coating.
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Fig. 5.6 Microstructure of as-fabricated coatings: SEM (BSE) images of: a) γ’ coating, b) γ/γ’
coating and c) β-NiPtAl coating; EBSD phase contrast map of the red box area in a - c: d) γ’
coating, e) γ/γ’ coating and f) β-NiPtAl coating; and corresponding color-coded inverse pole
figure (IPF) mapping g) - i) showing the different grain sizes of three coatings.
5.3.3 Isothermal oxidation performance of three Pt-diffused coatings
5.3.3.1 Elemental diffusion of three coatings
The elemental diffusion (Ni, Al and Pt) profiles of three as-fabricated coatings during
isothermal oxidation are shown in Fig. 5.7 - 5.9. For the Pt-diffused γ/γ’ coating, Pt showed
the same fluctuating trend as Al along the distance from the coating surface to the inner
superalloy, whereas Ni showed a contrary fluctuating trend (Fig. 5.8 a), indicating different
partition behaviours of these elements in γ and γ’ phase. Conversely, Pt (and Al) was uniformly
distributed throughout the Al-enriched γ’-phase coating (Fig. 5.7 a), which again confirmed the
pure γ’-phase microstructure. The Al distribution was relatively uniform throughout the β-
NiPtAl coating except for some small bumps (Fig. 5.9 a), which is attributed to the numerous
refractory precipitates (as also shown in Fig. 5.6 c). In summary, for the as-received coatings,
the average Al concentration: Pt-diffused γ/γ’ coating (~ 16.9 at. %) < the Al-enriched γ’-phase
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coating (~ 26.7 at. %) < β-NiPtAl coating (~ 41.8 at. %). While the Pt concentration of the β-
NiPtAl coating was considerably lower than that of other two coatings, which aligns with a
previous study [212].
Fig. 5.7 Ni, Pt and Al concentration evolution of the Al-enriched γ’ coating by EDX linescan
after a) 0 h, b) 20 h and c) 50 h oxidation.
After 20 h isothermal oxidation, the Pt and Al concentrations of the Al-enriched γ’-phase
coating near the surface only showed a little drop compared to that of as-fabricated coatings
(e.g. Pt concentration dropped from ~ 32 at. % to ~ 28 at. %, Fig. 5.7 a and b). On the other
hand, the Pt-diffused γ/γ’ coating exhibited a progressive reduction of both Pt and Al
concentrations near the surface, compared to the as-fabricated coating (Pt dropped from ~ 40
at. % to ~ 10 at. %, Fig. 5.8 a and b). Moreover, Pt concentration of the β-NiPtAl coating was
c)
a) b)
Ni
Al
Pt
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severely depleted after 20 h oxidation (Fig. 5.9 b) to an average concentration of ~ 5.0 at. %.
After 50 h oxidation, the β-NiPtAl coating showed the most severe Pt (and Al) depletion among
the three coatings. For example, Pt has diffused to a maximum depth of ~ 100 μm (the red
arrow in Fig. 5.9 c) for this coating. Conversely, for the Al-enriched γ’-phase coating, Pt only
diffused to a maximum depth of ~ 30 μm after 50 h (the red arrow in Fig. 5.7 c) and both Al
and Pt still remained high concentrations near the coating surface. In a word, the Pt and Al
interdiffusion with the substrate were much less pronounced for the Al-enriched γ’-phase
coating compared to the other coatings during isothermal oxidation.
Fig. 5.8 Ni, Pt and Al concentration evolution of the Pt-diffused γ/γ’ coating by EDX linescan:
a) 0 h, b) 20 h and c) 50 h oxidation.
b) a)
c)
Ni
Pt
Al
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Fig. 5.9 Ni, Pt and Al concentration evolution of the β-NiPtAl coating by EDX linescan: a) 0
h, b) 20 h and c) 50 h oxidation.
5.3.3.2 Oxide microstructure and growth kinetics
Fig. 5.10 shows the oxide phases (by XRD) on three coatings after different oxidation time at
1150 °C. Both α-Al2O3 and spinel (NiAl2O4) grew on the Pt-diffused γ/γ’ coating after 50 h
oxidation (Fig. 5.10 b). However, for the Al-enriched γ’-phase coating, spinel was only
identified after 100 h oxidation (Fig. 5.10 a). For the β-NiPtAl coating, only α-alumina was
detected in the oxide scale up to 100 h oxidation (Fig. 5.10 c). This demonstrates that although
the Al-enriched γ’-phase coating did not outperform the β-NiPtAl coating regarding the
selective oxidation of aluminium due to the much lower Al concentration (26.7 at. % compared
to 41.8 at. % for both as-fabricated coatings), it exhibited a better oxidation performance by
retarding Ni-oxide growth compared to the Pt-diffused γ/γ’ coating.
b)
c)
a)
Pt
Ni
Al
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Fig. 5.10 Glancing angle (3°) XRD patterns of the oxides on a) γ’ coating, b) γ/γ’ coating and c) β-NiPtAl coating after different oxidation time at 1150 °C.
a)
c)
b)
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Fig. 5.11 a) - c) cross-sectional SEM images of three coatings after 20 h oxidation at 1150 °C;
d) the magnified image of the red box area in c).
Fig. 5.11 shows the cross-sectional BSE images of the three coatings after 20 h isothermal
oxidation at 1150 °C. For the Al-enriched γ’-phase coating and the β-NiPtAl coating (Fig. 5.11
a and c), a pure α-Al2O3 layer was observed after 20 h oxidation, which agrees well with XRD
results (Fig. 5.10 a and c). A large number of alumina particles were observed near the
TGO/coating interface in the Al-enriched γ’-phase coating (red arrows in Fig. 5.11 a) because
of the internal oxidation. For the Pt-diffused γ/γ’ coating, local spinel formation above the α-
Al2O3 layer was observed, as shown by the red arrows in Fig. 5.11 b. It is noticed that XRD
did not detect any spinel in the oxide scale of the Pt-diffused γ/γ’ coating after 20 h oxidation
(Fig. 5.10 b), which is likely due to the small amount of spinel at this stage. Moreover, as can
be seen in Fig. 5.11 a, the Al-enriched γ’-phase coating remained γ’-phase structure near the
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coating surface after 20 h oxidation, as confirmed by EDX spectrum. The Pt-diffused γ/γ’
coating, on the other hand, exhibited a progressive γ’ to γ phase transformation near the
TGO/coating interface (Fig. 5.11 b), which is attributed to the Al depletion resulting from TGO
growth and interdiffusion with the substrate [185]. The β-NiPtAl coating maintained the β-
phase structure, but occasionally, γ’ precipitates can be seen at the grain boundaries of the β-
NiPtAl coating (red arrows in Fig. 5.11 d), in addition to some refractory metal precipitates
(black arrows in Fig. 5.11 d) [122].
Fig. 5.12 shows the TGO morphology of three coatings after isothermal oxidation for 50 h. At
this stage, some spinel formed locally on top of the continuous alumina layer on the Al-enriched
γ’-phase coating, as indicated by the red arrows in Fig. 5.12 a. And for this coating, no γ’ to γ
transformation took place near the TGO/bond coat interface. The TGO scale on the Pt diffused
γ/γ’ coating was essentially duplex in structure after 50 h (Fig. 5.12 b), with outer spinel layer
and inner α-Al2O3 layer. A thin layer consisting of predominantly γ-phase existed just below
the bond coat surface, which is due to γ’ to γ transformation as a result of Al depletion. For the
β-NiPtAl coating, the TGO scale still consisted of pure alumina after 50 h oxidation.
Furthermore, this coating maintained β-phase structure except for a small amount of γ’-phase
at the coating grain boundaries (Fig. 5.12 c). These observations were also consistent with XRD
results in Fig. 5.10.
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Fig. 5.12 Cross-sectional SEM backscattered electron (BSE) images of three coatings after 50
h oxidation at 1150 °C.
Fig. 5.13 Oxide thickness vs. isothermal oxidation time (at 1150 °C) for the three bond coat
systems.
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The TGO thickness vs. isothermal oxidation time are plotted in Fig. 5.13 for the three coatings.
The TGO on the Pt-diffused γ/γ’ coating grew much faster than the other two coatings up to
100 h oxidation time, which is attributable to the spinel formation of this coating. The Al-
enriched γ’-phase coating exhibited a slightly higher TGO growth rate, compared to the β-
NiPtAl coating due to the local spinel formation after 50 h oxidation (Fig. 5.12 a).
5.3.3.3 TGO spallation
Fig. 5.14 gives the optical surface images of three coatings after different isothermal oxidation
time at 1150 °C. The TGO scale remained intact for the Al-enriched γ’-phase coating and the
β-NiPtAl coating after 100 h oxidation (Fig. 5.14 a and c). Conversely, the oxide scale on the
Pt-diffused γ/γ’ coating exhibited noticeable spallation (bright areas in Fig. 5.14 b) after 50 h
oxidation. This indicates that the TGO spallation resistance of this new Al-enriched γ’-phase
coating is comparable to the conventional β-NiPtAl coating. Moreover, the Al-enriched γ’-
phase coating has exhibited significant improvement in terms of TGO lifetime compared to the
Pt-diffused γ/γ’ coating.
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Fig. 5.14 Optical surface images of three coatings after different isothermal oxidation time at
1150 °C.
5.3.4 Rumpling behaviour of three bond coats under cyclic oxidation
Fig. 5.15 - 5.17 illustrate the evolution of surface topography of three coatings under cyclic
oxidation, respectively. The surface topography recorded from an identical region of the Al-
enriched γ’-phase coating did not show significant changes from as-received condition up to
50 10-min cycles (Fig. 5.15 a - e). For detailed evaluation of the surface topography, the
roughness profiles along a line segment (the white line in Fig. 5.15 a) across the probed region
have been recorded at different stages of cyclic oxidation, as shown in Fig. 5.15 f. From the
statistical analysis, the roughness parameter for this line remained nearly constant during cyclic
oxidation: Rq (root mean square of roughness, calculated by Eq. 5.1) = 1.4 μm after 5 cycles
and Rq = 1.9 μm after 50 cycles. The small increase of Rq (~ 0.5 μm) was within the margins
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of measurement error. The Pt-diffused γ/γ’ coating exhibited similar behaviours, as illustrated
in Fig. 5.16. The surface topography for the as-received coating (Fig. 5.16 a) was almost the
same as that after 50 10-min cycles (Fig. 5.16 e). This is further confirmed by the roughness
profile evolution of a line segment (the white line in Fig. 5.16 a), as shown in Fig. 5.16 f where
Rq only slightly increased from 2.9 μm (5 cycles) to 3.4 μm (50 cycles). These investigations
indicate that both the Al-enriched γ’-phase coating and the Pt-diffused γ/γ’ coating did not
show detectable rumpling during the cyclic oxidation. Other researchers have also concluded
that Pt-diffused γ/γ’ coatings were resistant to the rumpling deformation during cyclic
oxidation [59, 213].
Conversely, the profilometer images recorded from an identical region of the β-NiPtAl coating
showed progressive changes during cycling test (Fig. 5.17). Clearly this coating has
demonstrated a tendency to roughen with thermal cycling. In addition, once rumpling was
initiated, the regions above the average surface persisted to bow up, whereas the regions below
the average surface continued to depress down with further cycling. This observation is in
agreement with the result reported by Chen et al. [118] on the NiCoCrAlY bond coat. Fig 5.17
f exhibits the roughness profile evolution of the white line in Fig. 5.17 a. These profiles clearly
showed that the corresponding positions along this line have moved either up or down
progressively with respect to the original surface plane. As a result, the roughness Rq has
increased from 0.9 μm after 5 cycles to 2.3 μm after 50 cycles. This suggests that the β-NiPtAl
coating showed significant rumpling increment during cyclic oxidation, which agrees with
some previous studies on β-NiPtAl bond coats [121, 206, 214].
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Fig. 5.15 Profilometer images of (a) as-fabricated Al-enriched γ’-phase coating surface and
after (b) 5 10-minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the
identical area; (f) the surface profiles and the corresponding Rq of the line shown in (a).
d
)
c
)
e 50 × 10-minute
10 × 10-minute 25 × 10-minute
f
a
)
b
)
5 × 10-minute As-received
Profile line in Fig. (f)
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Fig. 5.16 Profilometer images of (a) as-fabricated Pt-diffused γ/γ’ coating surface and after (b)
5 10-minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the identical area;
(f) the surface profiles of the line shown in (a).
As-fabricated
e
)
d
) c
)
b
)
a
)
f
50 × 10-minute
25 × 10-minute
5 × 10-minute
10 × 10-minute
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Fig. 5.17 Profilometer images of (a) as-fabricated β-NiPtAl coating surface and after (b) 5 10-
minute cycles, (c) 10 cycles, (d) 25 cycles and (e) 50 cycles recorded at the identical area; (f)
the surface profiles of the line shown in (a).
e
)
d
) c
)
50 × 10-minute
25 × 10-minute 10 × 10-minute
5 × 10-minute As-fabricated
f
a
) b
)
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5.4 Discussion
5.4.1 Pt and Al depletion of three coatings
The elemental diffusion profiles in Section 5.3.3.1 show that Pt depletion (due to interdiffusion
with the substrate) was more severe in the Pt-diffused γ/γ’ coating than that in the Al-enriched
γ’-phase coating during isothermal oxidation. In other words, Pt can remain relatively stable in
the pure γ’-phase microstructure. This is in agreement with Bai’s work [24], which has shown
that Pt is more stable in the γ’ phase than in the γ phase by thermodynamic calculations.
As for the Al depletion, firstly, the β-NiPtAl coating had the most severe Al depletion due to
interdiffusion with the substrate. While the Pt-diffused γ/γ’ coating exhibited a reduction of Al
depletion compared to the β-NiPtAl coating. Because the Pt-diffused γ/γ’ coating has the
chemical compatibility with the substrate that can mitigate the interdiffusion, as also reported
in [209]. Furthermore, the Al-enriched γ’-phase coating has experienced even less Al depletion
compared to the Pt-diffused γ/γ’ coating. This can be interpreted by the comparison of the local
interdiffusion flux of the Al-enriched γ’-phase coating/CMSX-4 diffusion couple and the Pt-
diffused γ/γ’ coating/CMSX-4 couple at specific positions during the oxidation. For simplic ity ,
we considered the Ni-Pt-Al ternary system, in which Ni was taken to be dependent and Al and
Pt were independent variables. Thus, the local interdiffusion flux of Al (𝐽𝐴𝑙) can be described
by Fick’s first law in terms of the two independent concentration gradients [54]:
𝐽𝐴𝑙 = −𝐷𝐴𝑙𝐴𝑙𝑁𝑖 𝜕𝐶𝐴𝑙
𝜕𝑥 − 𝐷𝐴𝑙𝑃𝑡
𝑁𝑖 𝜕𝐶𝑃𝑡
𝜕𝑥 (5.2)
where 𝐷𝐴𝑙𝐴𝑙𝑁𝑖 is the main-term interdiffusion coefficient of Al which relates the flux of Al to its
own concentration gradient; 𝐷𝐴𝑙𝑃𝑡𝑁𝑖 is the cross-term interdiffusion coefficient which accounts
for the chemical interaction between Pt and Al; 𝜕𝐶𝐴𝑙
𝜕𝑥 and
𝜕𝐶𝑃𝑡
𝜕𝑥 are the local concentration
gradients of Al and Pt, respectively.
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To calculate 𝐽𝐴𝑙 of the two couples at specific position 𝑥𝑖, the data for 𝐷𝐴𝑙𝐴𝑙𝑁𝑖 can be found in
[54] for different Ni-Pt-Al concentrations of the γ and γ’ phase, respectively. Specifically, 𝐷𝐴𝑙𝐴𝑙𝑁𝑖
for the Al-enriched γ’-phase coating/CMSX-4 diffusion couple can be found directly in [54],
while 𝐷𝐴𝑙𝐴𝑙𝑁𝑖 for the γ/γ’ coating was calculated according to the rule of mixture based on the
ratio of the γ and γ’ phase from the cross-sectional SEM image processing. As for the cross-
term interdiffusion coefficient, according to Hayashi et al. [54], there is no dependence of 𝐷𝐴𝑙𝑃𝑡𝑁𝑖
on the Pt and Al contents within the composition range studied (up to 25 at.% Pt addition in
the γ-Ni and γ’-Ni3Al alloys). Thus 𝐷𝐴𝑙𝑃𝑡𝑁𝑖 was taken as -2.4×10-10 and -0.9×10-10 cm2/s for the
γ and γ’ phase respectively in this study. 𝜕𝐶𝐴𝑙
𝜕𝑥 and
𝜕𝐶𝑃𝑡
𝜕𝑥 of the two couples can be calculated
from the elemental concentration profiles (Fig. 5.7 and 5.8). The measured concentration
profiles of Al and Pt were fitted using a cubic spline interpolation method, then 𝜕𝐶𝐴𝑙
𝜕𝑥 and
𝜕𝐶𝑃𝑡
𝜕𝑥
can be obtained at any 𝑥𝑖 of the diffusion couple. The interdiffusion coefficient in Eq. 5.2 has
the unit cm2/s, and the concentration has the unit mol/cm3. The elemental concentration with
the unit of atomic percent (at. %) in Fig. 5.7 and 5.8 can be converted into mol/cm3 by
introducing the molar volume (𝑉𝑚, unit is m3/mol). The average molar volume (𝑉𝑚) for the γ’
phase was estimated to be 6.83×10-6 m3/mol [215]. Since the molar volume varies very slightly
with the change of composition for the γ solid solution, the average 𝑉𝑚 for this phase was taken
as 6.68×10-6 m3/mol [215]. Then 𝑉𝑚 for the γ/γ’ coating was calculated based on the rule of
mixture. The local interdiffusion flux of Al (𝐽𝐴𝑙, calculated by Eq. 5.2) at specific positions
(𝑥𝑖=1, 5, 10 and 20 μm respectively) for these two couples after 20 h oxidation at 1150 ℃ is
compared in Fig. 5.18.
As can be seen from Fig. 5.18, the arrows indicate the direction of the local interdiffusion flux.
‘←’ represents the uphill diffusion from the inner part to the coating/TGO interface, in which
CHAPTER 5 THE AL-ENRICHED γ’-BASE BOND COAT FOR TBC APPLICATIONS
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the interdiffusion flux of Al (𝐽𝐴𝑙) has the opposite sign with the Al concentration gradient (𝜕𝐶𝐴𝑙
𝜕𝑥).
‘→’ represents the inward diffusion from the coating to the inner part of superalloy, in which
the interdiffusion flux has the same sign with the concentration gradient. Fig. 5.18 showed that
for 20 h oxidation, the Al-enriched γ’-phase coating exhibited the uphill Al diffusion at 𝑥𝑖=1,
5 and 10 μm, whereas the Pt-diffused γ/γ’ coating only exhibited uphill Al diffusion at 𝑥𝑖=1
μm. The cross-term interdiffusion coefficient 𝐷𝐴𝑙𝑃𝑡𝑁𝑖 was found to be negative in sign for both
coatings, suggesting that Pt has a negative chemical interaction with Al. In other words, Pt can
reduce the chemical activity of Al and promote the uphill diffusion of Al when the
concentration gradients of Al and Pt have opposite sign in the interdiffusion zone [216]. This
uphill diffusion can mitigate Al depletion during the oxidation process. Since the Al-enriched
γ’-phase coating exhibited uphill Al diffusion to a greater extent as shown in Fig. 5.18, the Al
depletion was less pronounced for this coating. This also coincided with the significant ly
reduced Al depletion of the γ’-phase coating compared to the γ/γ’ coating, as shown in Section
5.3.3.1.
The above calculations confirmed that the new γ’-phase coating exhibited less pronounced Al
depletion due to coating-substrate interdiffusion, which is beneficial to its lifetime. However,
there are a number of sources of error with this method to calculate 𝐽𝐴𝑙. For instance, the
concentration measurements by EDX are believed to be within the accuracy of ~ + 1.0 at. %
for each element. The cubic spline can reduce the concentration fluctuations measured by EDX.
But when the concentration gradient is really low, the fitting error can be relatively significant.
The assumption that the partial molar volume for each phase is concentration-independent
within the composition range studied can also be the error source. However, the trends and
agreements with experimental results are reasonable for the 𝐽𝐴𝑙 as shown in Fig. 5.18, which
provides an illustration for the elemental diffusion evolutions in Section 5.3.3.1.
CHAPTER 5 THE AL-ENRICHED γ’-BASE BOND COAT FOR TBC APPLICATIONS
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Fig. 5.18 The local interdiffusion flux of Al (𝐽𝐴𝑙) for the Al-enriched γ’-phase coating/CMSX-
4 diffusion couple and the Pt-diffused γ/γ’ coating/CMSX-4 diffusion couple at specific
positions (𝑥𝑖= 1, 5, 10 and 20 μm) after 20 h oxidation. The arrows represent the diffusion
direction. ←: uphill diffusion from the inner part to the TGO/coating interface; →: from the
coating to the inner part of superalloy.
5.4.2 Effect of bond coat composition on the selective oxidation of aluminium
The microstructural investigations (Section 5.3.3.2) confirmed that the oxide scale on the β-
NiPtAl coating was composed of exclusive alumina (without any Ni-oxides) up to 100 h
oxidation at 1150 °C. This is expected due to its high Al concentration (~ 41.8 at. %) which
ensures the selective oxidation of aluminium. The Al-enriched γ’-phase coating can promote
an exclusive growth of alumina and retard the Ni-oxide growth at initial oxidation compared
to the Pt-diffused γ/γ’ coating, although the Al concentration of the Al-enriched γ’-phase (~
26.7 at. %) is under the Al concentration minimum (~ 40.0 at. %) predicted in [24] for a
selective oxidation of aluminium. Two factors could be considered to explain this phenomenon.
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First, Chen et al. [55] proposed that large number of nanostructured grain boundaries in the
surface region of the NiCoCrAlY coating can provide numerous fast diffusion paths of
aluminium at initial stage of oxidation and therefore contributes to the establishment of an
exclusive alumina scale. The Al-enriched γ’-phase coating in this study, however, did not have
any nanostructured grains in the coating surface, as can be seen from the inverse pole figure
map of the as-fabricated coating (Fig. 5.6 a). This excluded this microstructural factor for
promoting the selective oxidation of aluminium. Second, Section 5.3.3.1 has shown that the Pt
concentration of the Al-enriched γ’-phase coating sustained at a relatively high level during
isothermal oxidation, especially at the coating surface. The Pt enrichment in the coating surface
can reduce the chemical activity of aluminium at the oxide/coating surface (Section 5.4.1),
which facilitates the aluminium diffusion flux from the inner part to the TGO/coating interface,
thus promoting the exclusive growth of alumina at initial stage of oxidation [68].
5.4.3 Rumpling behaviour of three coatings
Results in Section 5.3.4 suggest that both Pt-diffused γ/γ’ coatings and the Al-enriched γ’-phase
coatings showed negligible rumpling during cyclic oxidation, whereas the β-NiPtAl coating
experienced a significant higher degree of rumpling than the other two coatings. While the
substantial rumpling deformations of the β-NiPtAl coating are expected at this temperature
(1150 °C), the mechanisms responsible for the absence of rumpling in the new Al-enriched γ’-
phase coating are worth consideration. The following sections will discuss the possible
mechanisms by introducing a classical rumpling model firstly.
5.4.3.1 Balint and Hutchinson rumpling model
Rumpling in a TBC system is characterized by viable material and geometric parameters
including TGO thickening, lateral TGO growth strain, CTE mismatches, high temperature
strength of TGO, bond coat strength & phase transformation, etc. Balint and Hutchinson (B&H)
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166
[125] have proposed a comprehensive rumpling model that resolves these competing factors
and can be used to clarify the rumpling behaviour of coatings in experiments. In the B&H
model, the rumpling is driven by the lateral growth strain of TGO and occurs at a rate governed
by factors including the power-law creep of the bond coat and the plastic yielding of the TGO.
The power-law creep of the bond coat can be described by the following temperature-
dependent equation for the steady-state creep [204]:
휀̇ = 𝐶(𝜎
𝐸)𝑛exp (
−𝑄𝑐𝑟𝑒𝑒𝑝
𝑅𝑇) (5.3)
where 휀̇ is the creep rate; C is a constant; σ is the temperature-dependent equi-biaxial stress in
the coating; E is the temperature-dependent Young’s modulus; R is the gas constant and T is
the absolute temperature; 𝑄𝑐𝑟𝑒𝑒𝑝 is the activation energy for creep.
The equi-biaxial stress of the coating imposed by the coating/substrate CTE mismatch can
promote the coating creep, thus promoting the rumpling through the interaction between this
equi-biaxial stress in the coating and the normal traction imposed on the coating surface by
TGO [204]. At high temperatures, when the coating stress is relaxing, the undulation growth
of the coating is expected which is driven by the normal stress applied on the coating surface
by the compressed TGO layer. When the coating stress decays completely by creep, the
undulation growth is prohibited effectively. Because the cyclic oxidation scheme periodically
redefines the stress of the coating, the rumpling thus grows cycle-by-cycle [5].
5.4.3.2 B&H model applied to the Al-enriched γ’-phase coating
The first consideration regarding the absence of rumpling of the Al-enriched γ’-phase coating
compared to the β-NiPtAl coating is CTE mismatch. The γ’-phase coating has a smaller CTE
mismatch with the CMSX-4 substrate (CTE: ~ 15.0×10-6 °C-1 [217]) than that of the β-NiPtAl
coating [217]. This can result in the lower rumpling amplitude of the Al-enriched γ’-phase
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167
coating for a given thermal history according to the B&H model [125], because CTE mismatch
also contributes to the rumpling driving force.
The second consideration is the TGO growth rate. The B&H model has proposed that coatings
with thicker TGOs tend to have a faster rumpling rate because the thicker TGO imposes a larger
traction on the coating. However, the present experiments demonstrate that the Al-enriched γ’-
phase coating has a significantly lower rumpling rate in spite of its larger TGO thickness
(Section 5.3.3.2). This discrepancy can be explained by the improved creep strength of the γ’-
phase coating, as suggested by Jorgensen et al. [203]. In their simulation work, they have shown
that increasing the coating creep strength drastically inhibits rumpling even with thicker TGOs.
These simulations corroborate with the present study: the γ’-phase coating resists rumpling to
a large degree with only a slight increase in the total amplitude over the thicker TGO layer.
5.5 Summary
An Al-enriched γ’-phase coating was fabricated on the CMSX-4 substrate by the selective γ-
phase etching and a subsequent low temperature pack cementation aluminizing process. The
elemental diffusion profiles, TGO microstructure & spallation, and rumpling of the Al-enriched
γ’-phase coating have been investigated and compared to the industry-standard β-NiPtAl
coating and the Pt-diffused γ/γ’ coating. The following conclusions can be drawn:
1. The Al-enriched γ’-phase coating exhibited a comparable TGO spallation lifetime to
the β-NiPtAl coating and a significantly improved TGO lifetime compared to the Pt-
diffused γ/γ’ coating.
2. The Al-enriched γ’-phase coating exhibited a less pronounced coating-substrate
interdiffusion (Pt and Al depletion) compared to the other two coatings, which is
beneficial to the coating lifetime.
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3. The Al-enriched γ’-phase coating was more resistant to rumpling than the β-NiPtAl
coating, which mainly arose from the improved creep strength of the γ’-phase.
In summary, the new Al-enriched γ’-phase coating can resist high temperature rumpling
deformation while maintaining adequate oxidation properties and mitigating coating-substrate
interdiffusion, which is promising for TBC applications.
CHAPTER 6 CONCLUSIONS AND FUTURE WORK
169
Chapter 6 Conclusions and Future Work
6.1 Conclusions
Overall, this work has provided a new mechanistic understanding on the effects of the
underlying metalllic bond coat or substrate microstructure and chemical composition on the
TBC oxidation behaviour. It is concluded that the chemical composition of the underlying
coating or superalloy substrate has a direct effect on the oxidation behaviour of a TBC system,
from the early oxidation stage (Chapter 3) until the prolonged oxidation stage (Chapter 4 and
5). Specifically, the early stage oxidation of γ/γ’-based Ni-Al coatings with different Pt addition
(Chapter 3), the substrate composition effect on prolonged TBC lifetimes (Chapter 4) and a
new Pt diffusion bond coat design (Chapter 5) have been studied in this thesis. The main
conclusions are:
A. Pt addition has three effects on the early stage oxidation of γ/γ’-based Ni-Al coatings: 1)
retard θ-Al2O3 to α-Al2O3 transformation thus extending the transient θ-Al2O3 lifetime; 2)
promote the growth of Al2O3 and inhibit the growth of Ni-oxides; 3) significantly reduce
the TGO stress during the early stage of oxidation. Crystal orientation mapping results
show that the nucleation of α-Al2O3 is inhomogeneous along the oxide/coating interface
and might be related to the variation of coating compositions due to Pt additions. Spatially
resolved PLPS studies of the Ni-Pt-Al alloy with different Pt additions show that where Pt
stabilises the γ’ structure in nickel, the suppression of θ-Al2O3 to α-Al2O3 transition is
observed. Based on these findings, a new mechanism has been proposed to explain this Pt
effect on θ-Al2O3 to α-Al2O3 transformation: γ grains near the coating/oxide interface
promote the θ-Al2O3 to α-Al2O3 transformation while γ’ grains retard this transformation.
CHAPTER 6 CONCLUSIONS AND FUTURE WORK
170
B. Cyclic oxidation tests at 1200 °C have been carried out on TBCs with the CMSX-4 and the
René N5 substrate respectively to study the effect of substrate composition on TBC
lifetimes. It was found that TBCs based on the CMSX-4 superalloy had a 25% longer
average cyclic lifetime compared to that on the René N5 superalloy. Meanwhile, the failure
occurred mainly along the bond coat/TGO interface for TBC with the René N5 substrate,
whereas for TBC with the CMSX-4 substrate, a mixed failure path (along the bond
coat/TGO interface and within TGO) was observed. This indicated that the different
cracking behaviour at the TGO/bond coat interface for the two TBCs may originate from
the difference in the intrinsic interface toughness evolution. To confirm this, a strain-to-fail
method combined with the 3D-DIC technique was employed to measure the bond
coat/TGO interface toughness of TBCs with different substrates. Although the mode I
interfacial toughness (Γic) values were almost identical for the two substrates in the as-
deposited state (~ 30 J/m2), the interfacial toughness of René N5 specimens decreased faster
with oxidation time. This faster degradation of TGO/bond coat interfacial toughness for the
N5 substrate can be ascribed to the sulfur segregation at this interface.
C. An Al-enriched γ’-Ni3Al-base bond coat was successfully fabricated on the CMSX-4
substrate by the selective γ-phase etching and a subsequent low temperature pack
cementation aluminizing process. The elemental diffusion profiles, TGO microstructure &
spallation, and rumpling of this Al-enriched γ’-phase coating have been investigated and
compared to the industry-standard β-NiPtAl coating and the Pt-diffused γ/γ’ coating.
Compared to the two conventional diffusion coatings, the Al-enriched γ’-phase coating can
resist high temperature rumpling deformation while maintaining adequate oxidation
CHAPTER 6 CONCLUSIONS AND FUTURE WORK
171
properties and mitigating coating-substrate interdiffusion, which is promising for TBC
applications.
6.2 Future work
A. Study the early stage oxidation of NiCoCrAlY bond coats by scanning diffraction analysis
The findings in Chapter 3 indicated that the phase structure near the interface of the underlying
coating plays a key role in the transient Al2O3 to stable α-Al2O3 phase transformation rate
during early stage oxidation. This motivates further research by the scanning diffraction
technique to investigate the early stage oxidation behaviour of NiCoCrAlY bond coats with a
focus on the transient alumina transformation. Since the NiCoCrAlY bond coat is also two-
phase microstructure (β + γ), the alumina phase transformation mechanism relating to the
underlying coating phase structure might also be applied to this bond coat.
B. Study the effect of substrate composition on the lifetime of TBCs with β-NiPtAl bond coats
In Chapter 4, we have shown that substrate compositions can affect the lifetimes of TBCs with
the Pt-diffused γ/γ’ bond coats and illustrated the mechanisms behind this substrate effect.
However, the substrate composition effect on the TBCs with the β-NiPtAl bond coats is still
not fully understood. Further study is required to identify if the substrate composition plays a
role in the lifetime and failure mechanisms of this TBC system. Also, the interface toughness
measurement based on the strain-to-fail method combined with 3D-DIC technique can also be
utilised on this TBC system to replenish the current data base on this critical issue.
C. Prolonged cyclic oxidation tests on the Al-enriched γ’-phase coatings
Chapter 5 has demonstrated that the Al-enriched γ’-phase coating exhibited superior rumpling
resistance during short-term cyclic oxidation, indicating this new bond coat has high creep
strength. Thus, this new bond coat is expected to show promising cyclic oxidation performance
CHAPTER 6 CONCLUSIONS AND FUTURE WORK
172
compared to the β-NiPtAl bond coats and Pt-diffused γ/γ’ bond coats because of its
combination of rumpling resistance and improved Al content. In order to validate this and
further explore the application of this new bond coat at elevated temperatures, prolonged cyclic
oxidation tests (1-h holding time, 1150 ℃) are required. The cyclic oxidation lifetime of this
new bond coat can be compared to that of the conventional Pt diffusion coatings reported in
the literature. In addition, some high temperature mechanical properties (e. g. yield strength)
also have significant influence on the durability of coatings during service. To obtain a better
understanding of its high temperature oxidation behaviour, the evolution of mechanical
properties of the Al-enriched γ’-phase bond coat at high temperatures needs to be
experimentally determined.
REFERENCE
173
Reference
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