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Thermal Barrier Coatings – Durability Assessment and Life Prediction Robert Eriksson

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Thermal Barrier Coatings – Durability Assessment and Life Prediction Robert Eriksson
Linköping Studies in Science and Technology.
Dissertation No. 1527
Thermal Barrier Coatings
– Durability Assessment and Life Prediction
Robert Eriksson
Department of Management and Engineering
Linköping University, 581 83, Linköping, Sweden
http://www.liu.se
Linköping, August 2013
During the course of the research underlying this thesis, Robert Eriksson was
enrolled in the graduate school Agora Materiae, a doctoral program within the
field of advanced and functional materials at Linköping University, Sweden.
Cover:
Fractured plasma sprayed zirconia.
Printed by:
LiU-Tryck, Linköping, Sweden, 2013
ISBN 978-91-7519-569-8
ISSN 0345-7524
Distributed by:
Linköping University
Department of Management and Engineering
581 83, Linköping, Sweden
© 2013 Robert Eriksson
This document was prepared with LATEX, August 20, 2013
Abstract
Thermal barrier coating (TBC) systems are coating systems containing a metallic bond coat and a ceramic top coat. TBCs are used in gas turbines for thermal
insulation and oxidation resistance. Life prediction of TBCs is important as
high-temperature exposure degrades the coatings through mechanisms such
as thermal fatigue and the formation and growth of thermally grown oxides
(TGOs). This thesis presents research on durability assessment and life prediction of air plasma sprayed TBCs.
The adhesion of thermal barrier coatings subjected to isothermal oxidation,
thermal cycling fatigue and thermal shock was studied. The adhesion strength
and fracture characteristics were found to vary with heat treatment type.
The influence of interdiffusion between bond coat and substrate was studied on TBCs deposited on two different substrates. The thermal fatigue life was
found to differ between the two TBC systems. While fractography and nanoindentation revealed no differences between the TBC systems, the oxidation kinetics was found to differ for non-alumina oxides.
The influence of bond coat/top coat interface roughness on the thermal
fatigue life was studied; higher interface roughness promoted longer thermal
fatigue life. Different interface geometries were tried in finite element crack
growth simulations, and procedures for creating accurate interface models were
suggested.
The influence of water vapour and salt deposits on the oxidation/corrosion
of a NiCoCrAlY coating and a TBC were studied. Salt deposits gave thicker TGOs
and promoted an Y-rich phase. The effect of salt deposits was also evident for
TBC coated specimens.
A microstructure-based life model was developed using the Thermo-Calc
software. The model included coupled oxidation-diffusion, as well as diffusion
blocking due to the formation of internal oxides and pores. The model predicted Al-depletion in acceptable agreement with experimental observations.
iii
Preface
This thesis summarises the work I have done during my time as a Ph.D. student
at Linköping University, 2008–2013. The thesis consists of two parts. The second part, which is the main part of the thesis, consists of seven scientific papers
that summarises my work during the research project. The first part of the thesis gives background to the research and provides general information about
the topics studied more in detail in the appended papers. The first part also
gives the necessary knowledge for the non-specialist reader. The introductory
parts of this thesis is based on my licentiate thesis High-temperature degradation of plasma sprayed thermal barrier coating systems from 2011.
Many thanks to my supervisor, Sten Johansson, and others involved in the
research project, Håkan Brodin, Sören Sjöström, Xin-Hai Li and Lars Östergren,
for their help and support over the years.
I would also like to thank my many colleagues at the Division of Engineering
Materials for contributing to such a nice work environment full of creativity,
intellect, curiosity and humour. I would especially like to thank Kang Yuan with
whom I have had the opportunity to cooperate during my last years as a Ph.D.
student.
Robert Eriksson
v
List of papers
The thesis is based on the following papers:
I. R. Eriksson, H. Brodin, S. Johansson, L. Östergren, X.-H. Li, “Influence of
isothermal and cyclic heat treatments on the adhesion of plasma sprayed
thermal barrier coatings”, Surf. Coat. Technol., vol. 205, pp. 5422–5429,
2011.
II. R. Eriksson, H. Brodin, S. Johansson, L. Östergren, X.-H. Li, “Fractographic and microstructural study of isothermally and cyclically heat treated
thermal barrier coatings”, Surf. Coat. Technol., in press.
III. R. Eriksson, S. Johansson, H. Brodin, E. Broitman, L. Östergren, X.-H.
Li, “Influence of substrate material on the life of atmospheric plasma
sprayed thermal barrier coatings”, Surf. Coat. Technol., in press.
IV. R. Eriksson, S. Sjöström, H. Brodin, S. Johansson, L. Östergren, X.-H. Li,
“TBC bond coat-top coat interface roughness: influence on fatigue life
and modelling aspects”, To be published.
V. R. Eriksson, H. Brodin, S. Johansson, L. Östergren, X.-H. Li, “Cyclic hot
corrosion of thermal barrier coatings and overlay coatings”, Proceedings
of the ASME Turbo Expo 2013.
VI. K. Yuan, R. Eriksson, R. Lin Peng, X.-H. Li, S. Johansson, Y.-D. Wang, “Modeling of microstructural evolution and lifetime prediction of MCrAlY coatings on nickel based superalloys during high temperature oxidation”, Surf.
Coat. Technol., in press.
VII. R. Eriksson, K. Yuan, S. Johansson, R. Lin Peng, “Microstructure-based
life prediction of thermal barrier coatings”, Presented at MSMF7, 2013,
To appear in Key Engineering Materials.
vii
Acknowledgements
This research has been funded by the Swedish Energy Agency, Siemens Industrial Turbomachinery AB, GKN Aerospace Engine Systems, and the Royal Institute of Technology through the Swedish research programme TURBO POWER,
the support of which is gratefully acknowledged.
Kang Yuan, Mikael Segersäll, Jan Kanesund and Ru Lin Peng are acknowledged for contributing to Fig. 5, 6 and 7.
ix
Acronyms
APS
BC
BRT
CTE
CVD
EB-PVD
EBSD
EDS
FCT
FE
FEA
FEM
HVOF
InCF
MICF
PBR
PS
PVD
RE
SEM
TBC
TC
TCF
TCP
TET
TGO
VPS
WDS
Y-PSZ
YAG
YAP
air plasma spray
bond coat
burner rig test
coefficient of thermal expansion
chemical vapour deposition
electron beam physical vapour deposition
electron backscatter diffraction
energy dispersive spectroscopy
furnace cycle test
finite element
finite element analysis
finite element method
high-velocity oxy-fuel spray
intrinsic chemical failure
mechanically induced chemical failure
Pilling-Bedworth ratio
plasma spray
physical vapour deposition
reactive element
scanning electron microscope
thermal barrier coating
top coat
thermal cycling fatigue
topologically close-packed
turbine entry temperature
thermally grown oxide
vacuum plasma spray
wavelength dispersive spectroscopy
yttria partially stabilised zirconia
yttrium aluminium garnet
yttrium aluminium perovskite
xi
Contents
Abstract
iii
Preface
v
List of papers
vii
Acknowledgements
ix
Acronyms
xi
Contents
xiii
Part I Background and theory
1
1 Introduction
1.1 Background . . . . . . . . . . . . . . . . . . . . . . . . . . . .
1.1.1 Gas turbine development towards higher efficiency .
1.1.2 The importance of coatings . . . . . . . . . . . . . . .
1.2 Aim of this work . . . . . . . . . . . . . . . . . . . . . . . . . .
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2 Materials for high temperature applications
2.1 Physical metallurgy of systems containing Ni, Co, Fe, Cr and Al
2.1.1 Base materials . . . . . . . . . . . . . . . . . . . . . . . . .
2.1.2 Overlay coatings . . . . . . . . . . . . . . . . . . . . . . .
2.2 Thermal barrier coating systems . . . . . . . . . . . . . . . . . .
2.2.1 Top coat materials . . . . . . . . . . . . . . . . . . . . . .
2.2.2 Bond coat materials . . . . . . . . . . . . . . . . . . . . .
2.3 Manufacturing of TBCs . . . . . . . . . . . . . . . . . . . . . . .
2.3.1 Microstructure of thermal spray coatings . . . . . . . . .
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xiii
3 Oxidation of coatings
23
3.1 Formation of a protective oxide scale . . . . . . . . . . . . . . . . . 23
3.2 The reactive element effect . . . . . . . . . . . . . . . . . . . . . . . 28
3.3 Breakdown of the protective oxide scale . . . . . . . . . . . . . . . . 29
4 Thermal fatigue of coatings
4.1 Crack nucleation mechanisms . . . . . . . . . . . . . . . . . .
4.2 Crack growth mechanisms . . . . . . . . . . . . . . . . . . . .
4.3 Coating life assessment . . . . . . . . . . . . . . . . . . . . . .
4.3.1 Microstructure-based life models . . . . . . . . . . . .
4.3.2 The NASA model . . . . . . . . . . . . . . . . . . . . . .
4.3.3 Model suggested by Busso et al. . . . . . . . . . . . . .
4.3.4 Model suggested by Brodin, Jinnestrand and Sjöström
5 Experimental methods
5.1 Thermal fatigue . . . . . . . . . . . . .
5.2 Corrosion test . . . . . . . . . . . . . .
5.3 Adhesion test . . . . . . . . . . . . . .
5.4 Microscopy . . . . . . . . . . . . . . . .
5.4.1 Specimen preparation . . . . .
5.4.2 Scanning electron microscopy
5.5 Interface roughness measurement . .
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6 Discussion of appended papers
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Bibliography
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Part II Appended papers
67
Paper I: Influence of isothermal and cyclic heat treatments on the adhesion of plasma sprayed thermal barrier coatings
71
Paper II: Fractographic and microstructural study of isothermally and cyclically heat treated thermal barrier coatings
81
Paper III: Influence of substrate material on the life of atmospheric plasma sprayed thermal barrier coatings
93
Paper IV: TBC bond coat-top coat interface roughness: influence on fatigue life and modelling aspects
105
xiv
Paper V: Cyclic hot corrosion of thermal barrier coatings and overlay coatings
127
Paper VI: Modeling of microstructural evolution and lifetime prediction
of MCrAlY coatings on nickel based superalloys during high temperature oxidation
137
Paper VII: Microstructure-based life prediction of thermal barrier coatings
151
xv
Part I
Background and theory
1
Introduction
1.1 Background
The technology of gas turbines arose during the early to mid 1900s [1] and is
now commonly used for power production and aircraft propulsion. Gas turbines are closely integrated in today’s society and technology, and the development towards higher efficiency and fuel economy is naturally desirable [1–3].
Fig. 1 shows two examples of gas turbines: Fig. 1 a) shows a stationary gas
turbine for power production, and Fig. 1 b) shows an aircraft engine. Fig. 1 also
marks the major parts of a gas turbine: 1) the compressor, which compresses
the air, 2) the combustor, in which air and fuel are mixed and ignited, and 3)
the turbine which drives the compressor and provides the power output for, for
example, electric power production. The latter two, combustor and turbine,
operate in a very demanding high-temperature environment.
1.1.1 Gas turbine development towards higher efficiency
Since the performance of gas turbines depends on the temperature in the turbine part of the gas turbine [4], an increase in efficiency can be achieved by increasing combustion temperature [5–9]. Consequently, the development of gas
turbines has driven the service temperature to higher and higher levels. Fig. 2
shows an example from the aircraft industry where the turbine entry temperature (TET) has kept increasing since the 1940s.
Increasing operating temperatures offer several challenges in the field of engineering materials. As the operating temperature is driven to higher levels,
material issues such as oxidation, corrosion, creep and loss of strength are inevitable [4, 9–12]. The state-of-the-art metallic materials for high temperature
applications, the superalloys, are already operating at their maximum capacity
3
a)
b)
compressor
compressor
combustor
combustor
turbine
turbine
Figure 1: Gas turbines for power production and aircraft propulsion. a) Landbased gas turbine, SGT 750, for power production, (courtesy of Siemens Industrial Turbomachinery). b) Aircraft engine RM 12, used in JAS 39 Gripen, (courtesy of Volvo Aero Corporation).
and further increase in operating temperature cannot be achieved by further alloy development alone [4, 8, 10, 13–15]. This is illustrated in Fig. 2 which shows
the capability of typical superalloys compared to the TET; as seen, the TET now
exceeds the material capability.
Furthermore, the increasing demand for a more energy sustainable and environmental friendly society has drawn attention to the use of bio-fuels in gas
turbines [16]. The incorporation of bio-fuels in gas turbine technology may
cause changes in the operating conditions in the turbine which may also have
consequences for the metallic materials used there.
4
1600
Temperature, °C
1400
air cooling
thermal barrier coatings
TET
Material capability
1200
1000
800
1940
1950
1960
1970
1980
Year
1990
2000
2010
Figure 2: The increase in turbine entry temperature (TET) over time. The graph
also shows the temperature the material can withstand. Adapted from Reed [4].
The current and future research on gas turbine technology is therefore largely influenced by the striving for higher fuel efficiency, lower emissions and the
ability to use renewable fuels in gas turbines. The Swedish research programme
TURBO POWER, of which the research presented in this thesis is a part, aims
at achieving this. The programme is run as a collaboration between Siemens
Industrial Turbomachinery, GKN Aerospace Engine Systems, the Swedish Energy Agency and several Swedish universities. The research programme TURBO
POWER seeks to:
• Improve fuel efficiency of power-producing turbomachines, thereby reducing emissions and decreasing environmental strain.
• Improve fuel flexibility by making possible the use of alternative fuels.
• Reduce operating costs of power-producing turbomachines.
By developing technology and generating knowledge for university and industry, TURBO POWER will contribute to a more sustainable and efficient energy system in Sweden. The research aims at being applicable and governed by
needs.
1.1.2 The importance of coatings
As the temperature approaches the upper limit of material capability, phenomena such as creep, loss of mechanical properties, oxidation and corrosion occur
5
rapidly and limit the life of metallic materials [7, 9, 10, 13, 17–19]. As an example, Fig. 3 shows the variation of tensile strength with temperature for some
common superalloys. As seen in Fig. 3, superalloys cannot maintain their tensile strength at temperatures typical in gas turbine combustors and turbines;
the combustion temperature of gas turbines is even approaching the melting
temperatures of the base-elements in superalloys (nickel, cobalt and iron), see
Fig. 3.
1800
1600
Inconel 718
tensile strength, MPa
1400
Waspaloy
1200
1000
combustion temp.
Inconel 738
Inconel 939
Haynes 230
800
600
melting temp. of Ni, Co and Fe
Hastelloy X
400
200
0
precipitation hardened
solid-solution strengthened
0
200
400
600
800
1000
temperature, ◦ C
1200
1400
1600
Figure 3: Tensile strength of some superalloys as function of temperature.
By lowering the temperature below the point at which the alloys lose their
engineering properties, they can still be used as structural materials. The high
operating temperatures of today’s gas turbines – and the even higher temperatures of tomorrow’s gas turbines – are made possible by the use of air cooling
and thermal insulation in the form of thermal barrier coatings (TBC) [4, 8, 10,
13–15]. Air cooling, if too ample, has the disadvantage of reducing the achievable efficiency increase somewhat [13, 20] while thermal barrier coatings offer
an effective mean to provide insulation and oxidation resistance [5–8, 14, 18].
Fig. 4 a) shows a schematic drawing of a thermal barrier coating system;
the three parts of a thermal barrier system are: 1) substrate (component), 2)
bond coat (BC), and 3) top coat (TC); with time at high temperature, a layer
of thermally grown oxides (TGO) develops between the bond coat and the top
coat [7].
The top coat is made of a ceramic material with low heat conductivity and
thus provides the necessary thermal insulation. The metallic bond coat ensures good adhesion of the ceramic coating and provides oxidation resistance
[12, 17]. The effect of applying a TBC system onto a gas turbine component
is illustrated in Fig. 4 b): the top coat introduces a temperature gradient and
6
a)
hot combustion gases
oxygen
b) 1400
200
0
substrate
(component)
400
cooling air
substrate (component)
600
bond coat
thermally grown oxides
800
top coat
bond coat
temperature, ◦ C
1000
top coat
hot combustion gases
1200
distance from surface
Figure 4: Thermal barrier coating system. a) The parts of a thermal barrier coating system. b) The temperature through a coated component in a gas turbine.
Based on Stöver and Funke [5].
hence enables high combustion temperatures while avoiding high temperature
degradation of metallic parts.
The use of TBCs in gas turbines is, however, not entirely without its problems. Since the ceramic top coat and metallic bond coat have different coefficient of thermal expansion (CTE), stresses arise in the bond coat/top coat interface due to temperature variations (such as start and stop of the turbine).
Stresses are also introduced in the interface due to growth of the TGO layer.
The interface stresses eventually lead to failure of the TBC by spallation of the
top coat, which deprives the TBC system of its heat insulating capability.
1.2 Aim of this work
Thermal barrier coating systems currently offer an effective method for increasing gas turbine combustion temperature and thereby increasing efficiency [4,
8, 10, 13, 14]. To fully utilise protective coatings in gas turbines, reliable life prediction of TBCs must be achieved [5, 14, 17]. TBCs are only beneficial as long as
they adhere to the metallic parts which they are meant to protect. Understanding of the failure mechanisms of TBCs and the development of life models are
therefore important areas of research [4, 5, 12, 17].
The current research project has involved studies that contributed to the
understanding of: TBC durability, methods for evaluation of TBC life and durability, and life modelling aspects. The performed research has aimed at developing and improving life models for air plasma sprayed TBCs in gas turbines.
For this purpose, the research project has involved testing of several different
coating systems as well as the use of diverse testing methods.
7
Evaluation of durability of TBCs has included: isothermal oxidation, furnace cycling, burner rig test, corrosion test, adhesion tests on thermally degraded specimens, nanoindentation as well as microscopy studies on microstructure and oxide composition and growth kinetics. The study includes the
investigation of several degrading mechanisms of TBCs during isothermal and
cyclic high-temperature exposure: fatigue damage, interface TGO growth, influence of substrate material on life, influence of BC/TC interface roughness on
life, and cracking and sintering of the top coat. Life prediction was tried both
from a fracture mechanics point of view and from an oxidation/interdiffusionbased point of view.
8
2
Materials for high temperature
applications
High temperature materials are materials that can operate at temperatures close to their melting temperatures while still maintaining many of the typical
room temperature characteristics of engineering materials, such as high strength and microstructural stability [4, 10, 11].
The base material makes up the structural parts of the gas turbine and their
chemistry may often be chosen for good mechanical properties rather than resistance to environmental degradation, i.e. oxidation and corrosion [11]. Three
classes of alloys: Ni-base, Co-base and Fe–Ni-base, collectively referred to as
superalloys, have shown to have good to excellent high temperature properties and are widely used as base material for high temperature applications
[4, 10, 11].
2.1 Physical metallurgy of systems containing Ni, Co,
Fe, Cr and Al
2.1.1 Base materials
In superalloys, the solid-solution γ-Ni phase – which has the face centred cubic
(FCC) atomic arrangement – constitutes the matrix phase. Ni-base superalloys
can be solid-solution strengthened, such as Haynes 230 and Hastelloy X, or precipitation hardened, such as Waspaloy and Inconel 738, 939 and 718. As seen
in Fig 3, precipitation strengthened alloys typically have higher strength than
solid-solution strengthen alloys and are used in more demanding high temperature environments [11]. Solid-solution strengthened materials have ad9
vantages when it comes to processing and have, for example, better weldability [11]; they can also be manufactured in complex geometries from powders
through laser melting techniques, see for example Saarimäki [21].
For solid-solution strengthened alloys, the alloying elements are chosen from Fe, Co, Cr, Mo, W, Ti and Al [10]. Al, Cr, W and Mo are potent solid-solution
strengtheners largely due to their different atomic radius compared to Ni [10].
For precipitation hardened alloys the alloying elements are typically chosen
from: Al, Ti, Ta, and sometimes Nb which promotes the formation of the γ0
or γ00 precipitates in the γ-matrix [4, 10, 11], shown in Fig. 5. In addition, minor
amounts of elements like Hf, Re, Zr, C and B may be added for various purposes.
The compositions of some common Ni-base alloys are given in Table 1.
a)
b)
TCP
γ
γ
γ
γ/γ
2 µm
2 µm
Figure 5: Some microconstituents in Ni-base alloys. a) Inconel 792 showing γ0
precipitates in a matrix of γ with secondary γ0 . b) CMSX-4 showing γ0 precipitates in a γ-matrix; some TCF phases can also be seen.
Table 1: Composition of some Ni-base alloys
Alloy
Haynes 230
Hastelloy X
Inconel 738
Inconel 939
Inconel 718
Waspaloy
a
b
Ni
Co
Fe
Cr
W
Mo
Al
Ti
Nb
Ta
Si
C
B
57a
47a
61.4a
47.3a
52.5
58a
5b
1.5
8.5
19
1b
13.5
3b
18
–
0.5b
18.4a
2b
22
22
16
22.5
19
19
14
0.6
2.6
2
–
–
2
9
1.75
–
3.1
4.3
0.3
–
3.4
1.9
0.5
1.5
–
–
3.4
3.7
0.9
3
–
–
0.9
1
5.1
–
–
–
1.75
1.4
–
–
0.4
1b
–
0.2b
0.35b
0.15b
0.1
0.1
0.17
0.15
0.08b
0.08
0.015b
0.008b
0.01
0.01
0.006b
0.006
balance
maximum
The γ0 phase is an aluminide with formula Ni3 (Al, Ti); the Al and Ti may be
substituted by Ta and Nb, and the Ni can, to some extent, be substituted by Co
10
or Fe [11]. The γ0 phase is an ordered phase with the L12 structure. The morphology of the γ0 precipitates depends on their mismatch with the surrounding parent lattice and includes: cubical, small spherical particles and arrays
of cubes [4, 11]. Modern precipitation hardened alloys may contain & 60 % γ0
[4, 11]. An interesting characteristic of γ0 is its increasing tensile strength with
increasing temperature [11].
For Ni–Fe alloys, such as Inconel 718, the addition of Nb may cause the precipitation of γ00 -Ni3 Nb [10, 11] which acts as the primary strengthening microconstituent. Alloys that rely on the strengthening from γ00 -Ni3 Nb are limited
to operating temperatures below ∼ 650 °C as the tetragonal γ00 -Ni3 Nb otherwise will transform to a stable orthorhombic δ-Ni3 Nb which does not add to
strength [11].
The addition of C and B enables the formation of carbides and borides.
Carbide formers include Cr, Mo, W, Nb, Ti, Ta and Hf, which form carbides of
various stoichiometry, such as MC, M23 C6 and M6 C. Common boride formers
are Cr and Mo, which form M3 B2 ; boron tends to segregate to grain boundaries [10, 11]. Carbon and boron performs an important role as grain boundary strengtheners and are consequently added in greater amounts to polycrystalline alloys [4].
MC carbides typically form at high temperatures, for example during solidification and cooling in the manufacturing process, while M23 C6 and M6 C form
at lower temperatures: 750–1000 °C [10]. The MC carbide typically forms from
Ti, Hf and Ta [4, 11], while the M23 C6 is promoted by high Cr contents and the
M6 C is promoted by large fractions of W and Mo [10]. The MC carbide may form
within grains as well as at grain boundaries; the M23 C6 carbides are preferably
formed at grain boundaries.
The MC carbides may decompose to form M23 C6 and M6 C carbides during
high temperature exposure during operation or heat treatment [22, 23]. The
following reactions have been suggested [10]:
MC + γ
M23 C6 + γ0
(A)
MC + γ
M6 C + γ0
(B)
A group of intermetallics generally considered harmful to Ni-base alloys,
are the topologically close-packed (TCP) phases, such as the σ phase. These
phases may form in alloys rich in Cr, Mo and W [4]. The σ phase has the general
formula (Cr, Mo)x (Ni, Co)y [10]; it may have a plate or needle-like morphology
and may appear in grain boundaries, sometimes nucleated from grain boundary carbides [10, 11].
11
2.1.2 Overlay coatings
Overlay coatings are deposited on top of the substrate without interacting much
with the substrate [13, 24]. This can be contrasted to diffusion coatings which
are coatings that are formed through interdiffusion with the substrate and the
coating is formed as the coating elements interact with the substrate [13, 24].
Overlay coatings are deposited by methods such as plasma spraying (PS),
electron beam physical deposition (EB-PVD) and high velocity oxy-fuel spraying (HVOF). These deposition methods use alloy feedstocks and the deposited
coating may therefore have a composition completely different from the substrate. Overlay coatings are often chosen from the MCrAlX family of alloys
where M is either Ni, Co or Fe or a combination of them; X denotes additions
of reactive elements (RE), which in various ways improve the properties of the
coating.
As for the base material, MCrAlX coatings consist of a γ-matrix with the
aluminium largely bound in aluminides. The γ0 aluminide may be present in
the microstructure, but for such large amounts of Al as are commonly used in
MCrAlX alloys, another aluminide forms: β-NiAl [15, 25]. For MCrAlX coatings, most of the aluminium is bound in this phase and the two main microconstituents of many MCrAlX coatings are γ and β. In addition, MCrAlX may
contain chromium rich σ−(Cr, Co) and α-Cr [25]. The latter may occasionally
precipitate in the β phase [25, 26]. Thus, a typical MCrAlX alloy may have microstructures such as: γ + β or γ/γ0 + β/α both with the possible addition of
σ−(Cr, Co) [12, 27–29].
Fig. 6 shows the phases present at high temperature for the Ni–Cr–Al system with different additions of Co. Fig. 7 and 8 shows the microstructure of two
MCrAlX coatings which have been cooled in air from high temperature; the figures show how electron backscatter diffraction (EBSD) and energy dispersive
spectroscopy (EDS) can be used to identify the γ, γ0 , β and σ phases.
2.2 Thermal barrier coating systems
A protective coating for high temperature applications must lower the temperature of the substrate and provide the oxidation and corrosion resistance which
the base material lacks. The requirements on a protective coating can be summarised as [13]:
• The coating must have low thermal conductivity.
• The coating needs to have good oxidation and corrosion resistance.
12
• A protective coating must have high melting temperature and retain its
structural integrity in the full interval of operating temperatures.
• The coating should have a coefficient of thermal expansion close to the
substrate on which it is deposited to avoid thermal mismatch.
As no single material possesses all of those properties, protection of superalloys is typically achieved by material systems containing an insulating coating
(top coat) deposited on top of an oxidation resistant coating (bond coat). The
bond coat also provides adhesion of the top coat to the substrate. A TBC system
is shown in Fig. 9.
a)
b)
30
T=1100 oC
10 wt.% Co
25
20
γ'
Cr
15
γ
10
5
γ' + β
γ'
γ+
γ'
γ+
wt.
%
wt.
%
β+σ
+β
10
γ+β+σ
5
β
0
5
10
15
20
25
30
0
5
10
15
wt.% Al
c)
d)
T=1100 oC
20 wt.% Co
25
25
Cr
wt.% Al
γ'
γ'
5
β
0
20
25
30
0
5
10
γ'
β
γ' + β
γ+
γ'
γ+
γ'
15
γ
10
'+β
10
γ+β
15
+β
γ+γ
β+σ
γ' + β
5
wt.
%
Cr
%
wt.
γ+β
γ
0
0
30
20
γ+β+σ
10
5
25
30
20
15
20
wt.% Al
30
T=950 oC
20 wt.% Co
β
γ' + β
γ'
0
0
β
β
γ+
γ' +
γ+
γ
γ+β
Cr
20
γ + γ'
15
25
γ+
T=950 oC
10 wt.% Co
30
15
20
25
30
wt.% Al
Figure 6: Phase diagrams for some NiCoCrAl alloys established by ThermoCalc. a) NiCrAl + 10 wt.% Co at 950 °C b) NiCrAl + 10 wt.% Co at 1100 °C c) NiCrAl
+ 20 wt.% Co at 950 °C d) NiCrAl + 20 wt.% Co at 1100 °C
13
a)
b)
γ
BCC
β
FCC
γ
5 µm
c)
d)
Cr
Al
Figure 7: Microstructure in a NiCoCrAlY coating analysed by EBSD and EDS: a)
electron micrograph, b) EBSD results showing crystal structure, c) EDS results
showing Al content, and d) EDS results showing Cr content.
a)
b)
β
σ
β
γ
γ
σ
50 µm
Figure 8: Microstructure in a NiCoCrAlY coating analysed by EBSD: a) electron
micrograph, b) phases identified by EBSD.
14
TC
TGO
BC
substrate
100 μm
Figure 9: The components in a thermal barrier system: substrate, bond coat
(BC), thermally grown oxides (TGO) and top coat (TC).
2.2.1 Top coat materials
The top coat is the part of the TBC system that provides thermal insulation. As
any insulation, the top coat must be combined with internal cooling to keep
the temperature low in the substrate; the top coat only introduces a steep temperature gradient. The temperature drop in a top coat, 300 µm in thickness, can
be as high as 200–250 °C [4, 7, 11, 13].
The 6–8 wt.% yttria partially-stabilised zirconia (Y-PSZ) has arisen as the industry standard for top coat material [30]. This is largely due to its combination
of low thermal conductivity and relatively high coefficient of thermal expansion
[7, 30]. The software CES provides a convenient tool for illustrating this; Fig. 10
shows a diagram of thermal conductivity and thermal expansion for a number
of technical ceramics and some Ni-based alloys; it can be seen that zirconia has
the desired combination of low thermal conductivity and high thermal expansion.
Pure zirconia (ZrO2 ) is allotropic: monoclinic up to 1170 °C, tetragonal in
the interval 1170–2370 °C and cubic up to the melting point at 2690 °C. The tetragonal to monoclinic transformation is martensitic in nature and involves a
3–5 % volume increase that induces internal stresses which compromise the
structural integrity of the ceramic [13, 31]. The tetragonal–monoclinic transformation is problematic since it occurs in the range of the operating temperatures in gas turbines.
The detrimental phase transformation can be avoided by stabilising the tetragonal phase. Various oxides, such as CaO, MgO, Y2 O3 , CeO2 , Sc2 O3 and In2 O3
15
Thermal conductivity (W/m.°C)
100
alumina
beryllia
boron carbide/nitride
graphite
magnesia
mullite
sapphire
silica
silicon carbide/nitride
tungsten carbide
sialon
10
nickel alloys
zirconia
1
0.2
0.5
1
2
5
10
Thermal expansion coefficient (µstrain/°C)
20
Figure 10: Thermal conductivity and thermal expansion coefficient for zirconia
compared to some other ceramics and Ni-alloys. Chart from CES EduPack 2012,
Granta Design Limited, Cambridge, UK, 2012.
[13, 15, 30], can be added to stabilise the tetragonal phase, but yttria has become the most common. The optimum amount of 6–8 wt.% of yttria is based
on the work of Stecura [32] who showed that TBCs with ∼ 6 wt.% had the highest
fatigue life in a thermal cycling test; see Fig. 11.
The phase being stabilised by the addition of 6–8 wt.% Y2 O3 is the nontransformable tetragonal phase, t0 , which is stable from room temperature to
approximately 1200 °C [4, 7, 30]. The t0 phase is formed by rapid cooling during coating deposition and is a metastable phase [30]. At high-temperature exposure, the t0 phase starts to transform to the equilibrium tetragonal and cubic phases. The t0
cubic + tetragonal transformation occurs as the Y-PSZ is
only partially stabilised. The addition of & 11 wt.% Y2 O3 would stabilise the
cubic phase from room temperature to the melting temperature and thus enable higher operating temperatures, but, as shown by Fig. 11, that would give
a shorter fatigue life. The high-temperature transformation of t0
cubic +
monoclinic transformation on
tetragonal enables the undesired tetragonal
cooling [33]. Hence, there exists an upper limit to the practical operating temperature of partially-stabilised zirconia.
16
400
350
Cycles to failure
300
250
200
150
100
50
0
5
10
15
wt.% Y2 O3
20
25
Figure 11: The thermal fatigue life of yttria-stabilised zirconia as function of
yttria content. Adapted from Stecura [32]
2.2.2 Bond coat materials
While the Y-PSZ top coat provides the necessary thermal insulation, it does not
offer any protection against oxidation and corrosion. The Y-PSZ coating readily
lets oxygen through and causes the underlying metal to oxidise [30]. To prevent
the substrate from oxidising, an oxidation resistant bond coat is incorporated
between the substrate and the top coat. The bond coat is chosen from alloys
with excellent oxidation properties. Furthermore, the bond coat improves adhesion between the top coat and the substrate, particularly for plasma sprayed
coatings.
While bond coats can be made from diffusion coatings [34], overlay coatings
are probably the most used and most developed for use as bond coats. Overlay
coatings enable elaborate alloy design as overlay coatings are independent in
chemistry from the substrate on which they are deposited; therefore, numerous
variations on the MCrAlX concept exist: Ni–(0–30 wt.% Co)–(10–30 wt.% Cr)–
(5–20 wt.% Al)–(. 1 wt.% Y) covers the range of many bond coat compositions.
Bond coat alloys contain the addition of one or several reactive elements. The
purpose of the REs is often to improve oxide scale adhesion; even RE additions
in the order of ∼ 0.1 wt.% may increases adhesion of the Al oxide scale [35]. Y
is the most widely used (typically . 1 wt.%) [7, 9, 12, 17, 30, 36, 37], but other
common additions include: Ce, Hf, Zr, Si, La, Re and Ta [13, 22, 30, 36].
MCrAlY coatings achieve oxidation and corrosion resistance through the
17
formation of a protective oxide scale in the bond coat/top coat interface. Such
protective scales need to be: stable at high temperatures, dense, slow-growing
and exhibit good adhesion to the coating [15]. Three oxides, alumina (Al2 O3 ),
chromia (Cr2 O3 ) and silica (SiO2 ), have the potential to fulfil these requirements [15, 38]. At temperatures common in gas turbines, a continuous layer
of alumina is usually the most beneficial for TBC life [11, 12].
The necessary ability to form a layer of protective alumina influences the
choice of chemistry for these alloys. The interfacial TGO is protective only as
long as it consists of predominantly Al2 O3 , and as long as it is intact and adherent to the bond coat. The chemistry of the bond coat must be chosen to assure
that: 1) aluminium is the preferred oxidising species, 2) the alumina has good
adherence to the bond coat and 3) the alumina is reformed if it is damaged.
A prediction of what kind of oxides an alloy will form can be obtained by an
oxide map, such as the one showed in Fig. 12 for the Ni–Cr–Al system. There
exists a critical Al content below which alumina cannot be formed. For example, Fig. 12 shows that ∼ 20 wt.% Al, (∼ 35 at.% Al), is needed to ensure Al2 O3
growth in a Ni–Al system. However, the addition of Cr promotes the formation
of a protective Al2 O3 scale [15]; with the addition of 5 wt.% Cr, (∼ 5 at.%), the
alloy can form Al2 O3 at an Al content as low as ∼ 5 wt.%, (∼ 10 at.%).
0 40
10
at.
%
Al2O3
20
Al
20
%
at.
Cr
30
30
10
Cr2O3
40
60
70
NiO
80
at.% Ni
90
0
100
Figure 12: Oxide map for the Ni–Cr–Al system at 1000 °C. Areas denoted Cr2 O3
and NiO may also give internal oxidation of Al2 O3 . Based on Wallwork and Hed
[39].
Al and Cr are consequently added in amounts of ≥ 5 wt.% to improve oxidation and corrosion resistance by assuring the formation of a protective alumina
scale. The composition of the bond coat must also be chosen to account for the
18
depletion of aluminium during high temperature exposure by consumption of
Al through oxidation and interdiffusion with the substrate; most bond coats
are consequently quite rich in Al [30]. As the Al content in the coating drops,
the β and γ0 phases will dissolve [22]. Two possible decomposition routes are
[11, 15, 26, 40]:
β
γ
β
γ + γ0
(C)
γ
(D)
2.3 Manufacturing of TBCs
TBC systems are manufactured by methods belonging to process families such
as thermal spraying, physical vapour deposition (PVD) and chemical vapour
deposition (CVD). The group of manufacturing methods collectively referred
to as thermal spraying includes processes such as plasma spraying and highvelocity oxy-fuel spraying, both commonly used for manufacturing of TBC systems [5, 7, 14, 19, 30]. Plasma spraying can be conducted in air or in vacuum and is, accordingly, referred to as atmospheric plasma spraying (APS) and
vacuum plasma spraying (VPS) or, alternatively, low pressure plasma spraying
(LPPS).
The raw materials for manufacturing of bond coats and top coats are typically in powder form. The plasma spray process uses a plasma jet to melt the
feedstock powder into droplets which are sprayed onto the substrate: powder
is introduced by a carrier gas into the plasma jet, melted and propelled towards
the substrate [19]. The characteristics of plasma sprayed coatings are largely
influenced by spraying conditions such as plasma jet velocity and the droplet
dwell time in the plasma jet [19].
plasma gas cathode powder inlet
anode
cooling water
plasma flame
spray stream
Figure 13: Schematic drawing of a plasma gun. Based on Ref. [41].
A schematic drawing of a plasma gun is shown in Fig. 13. The plasma gas, for
example argon, is brought into the plasma gun and led through an electric field
19
that ionises the gas to produce plasma; the plasma may reach temperatures as
high as 20 000 °C [19]. Due to the high temperature, the anode is water cooled
and the cathode is typically made from tungsten which has a sufficiently high
melting temperature and is a good thermionic emitter [19].
2.3.1 Microstructure of thermal spray coatings
The plasma spraying process gives rise to a very characteristic microstructure
where droplets from the spraying process can be discerned as flat, so called,
splats. As the molten droplets impact on the substrate, they form thin discshaped splats which cool on impact and solidify rapidly; for metal coatings,
with a speed of up to 106 K/s [19].
Atmospheric plasma sprayed metallic coatings have microstructures that
include constituents such as splats, oxide inclusions/stringers, pores and unmelted or partially melted particles. The microstructural characteristics of an
APS deposited bond coat are shown in Fig. 14 a) and can be contrasted to a VPS
deposited bond coat, shown in Fig. 14 b), whose characteristic features are the
absence of oxide stringers and lower porosity. The lower oxide fraction in VPS
coatings are caused by the spraying being performed in vacuum. The HVOF
process produces metallic coatings similar in appearance to the VPS coating.
For ceramic coatings, the rapid solidification typically causes a columnar
grain structure within each splat [42, 43], shown in Fig. 15 a). The typical splaton-splat structure is easily seen in Fig. 15 b) where it can also be seen that the
splats segment by forming a cracked-mud-like pattern of microcracks. Such
cracking is due to stresses caused by contraction during the rapid cooling of
the splat [42]. Fig. 15 b) also shows interlamellar delaminations [30, 42, 44–
46]. These crack-like voids are caused by the low area of contact between splats
which may be as low as 20 %[47]. Both the splat microcracks and interlamellar
delaminations can be readily seen on cross-sections, as shown in Fig. 15 c).
The APS process also gives rise to porosity, see Fig. 15 d). In the top coat, such
porosity is desirable as it decreases the thermal conductivity of the coating [30].
Porosity levels in TBCs typically lie in the interval 5–20 %.
20
a)
b)
TC
BC
A
TC
BC
A
A
substrate
50 μm
substrate
100 μm
Figure 14: Microstructure of plasma sprayed MCrAlY coatings. a) APS coating showing oxide inclusions/stringers, arrows, and partially melted particles,
marked by A. b) VPS coating without oxide stringers and with lower porosity.
a)
b)
C
A
B
1 μm
c)
5 μm
d)
B
D
A
5 μm
50 μm
Figure 15: Microstructural characteristics of an APS top coat. a) Columnar grain
structure in a splat. b) Fractured top coat showing the splat-on-splat structure,
C; interlamellar delaminations, B; and cracking of the splats, A. c) Cross-section
of a top coat showing interlamellar delaminations, B, and through-splat cracks,
A. d) Cross-sectioned top coat showing porosity, D.
21
3
Oxidation of coatings
The oxidation kinetics of high temperature alloys typically obey an Arrheniustype equation [48]
Q
k = k 0 e − RT
(1)
where k is the oxide growth rate constant, Q is the activation energy, T the temperature in K, R = 8.314 J/(mol K) is the gas constant and k 0 is a constant. Oxidation rate consequently increases exponentially with temperature and oxidation at high temperatures may be very fast. The formation of a protective layer
of BC/TC interface TGOs is essential for oxidation resistance but is a sacrificial
process during which the coating is consumed. Oxidation is a degrading mechanism that will eventually lead to the breakdown of the protective TGOs and
might induce failure of TBCs.
The oxidation of the BC can be divided into three stages, shown in Fig. 16:
1) a transient stage of simultaneous oxidation of all oxide-forming species in
the bond coat, 2) a steady-state stage of formation and growth of a protective
oxide scale, and 3) a breakaway stage of rapid oxidation and spallation [38].
The second stage gives rise to a protective oxide scale whereas the third stage
causes failure of the TBC system.
3.1 Formation of a protective oxide scale
The transient stage is the stage of oxidation before a continuous oxide layer has
formed on the metal surface and during which all oxide-forming species in the
alloy (Ni, Co, Cr, Al) might form oxides. The composition of the transient oxides is influenced by parameters such as: temperature, partial oxygen pressure,
23
oxide scale thickness
breakaway
transient
steady-state
high temperature exposure time
Figure 16: Schematic drawing of the three stages of oxidation: short stage of
transient oxidation, steady-state oxidation, and breakaway oxidation. Based on
Hindam and Whittle [38].
coating composition and coating microstructure [49]; low partial oxygen pressure, for example, may promote the formation of alumina [50]. Each element’s
affinity for oxygen will determine how it forms oxides. An Ellingham diagram,
such as the one in Fig. 17, provides information at which oxygen partial pressure an oxide can form according to [51]:
M /M x O y
pO2
= exp
∆G ◦
RT
(2)
This is the dissociation pressure of the oxide; the partial pressure of oxygen
must be higher than the dissociation pressure for the oxide to form. An Ellingham diagram can be used to rank the order in which the oxides will form; as
seen in Fig. 17, Al will oxidise at lower oxygen partial pressure and thus oxidises
more easily than, for example, Ni and Co.
The transient stage is usually quite short, typically . 1 h for Ni–Cr–Al systems oxidised at 1000–1200 °C [52, 53]. Transient oxides include Cr2 O3 , NiO,
CoO, spinel type (Ni, Co)(Cr, Al)2 O4 and various forms of alumina: γ-, θ-, αAl2 O3 [40, 49, 52–55].
The transition from transient oxidation to the slower steady-state stage of
growth occurs when a continuous oxide layer is formed and the oxidation rate
becomes controlled by the diffusion rate of oxygen and metal ions through
the oxide layer. Such diffusion controlled oxidation is typically described by
a power-law expression:
1
h TGO = h 0 + kt n
(3)
where h TGO is the thickness, (or weight gain per area), of the formed oxide, h 0 is
the thickness of the transient oxides, k is the growth rate constant and t is the
24
0
Free energy of formation ΔG◦ , kJ
-200
+
2 Ni
2 Co +
-400
4
3
-600
Cr +
O2 =
O
Si + 2
2
3
O2 =
2 NiO
CoO 2
O2 =
Cr 2O 3
= SiO 2
-800
4 Al
3
+ O2
=
2 Al 2O 3
3
-1000
-1200
400
600
800
1000 1200
Temperature, K
1400
1600
1800
Figure 17: Ellingham diagram showing the free energy of formation of common
oxides in coatings.
high temperature exposure time. The classical oxidation law is parabolic (n = 2)
[56] but subparabolic models (1/n < 0.5) are also in use [12, 57–59]; particularly
the cubic law (n = 3) has become common.
Protective oxide scales can be provided by Al, Cr and Si which form Al2 O3 ,
Cr2 O3 and SiO2 [11, 12, 38]. At high temperature, Al2 O3 is usually the protective oxide. The use of Cr2 O3 -forming coatings is restricted to somewhat lower
temperatures, (. 950 °C [11, 15]), as Cr2 O3 may decompose to volatile CrO3 and
evaporate according to [11, 12, 60]:
Cr2 O3 (solid) + 32 O2 (gas)
2 CrO3 (gas)
(E)
The use of SiO2 -forming coatings is also limited to lower temperatures as they
may form low-melting or brittle phases [15].
A protective layer of interface Al2 O3 can be seen in Fig. 18; Fig. 18 a) shows a
fracture surface produced by tearing off the top coat, thus exposing the underlying interface TGO, and Fig. 18 b) shows a polished cross-section of a layer of
interfacial TGO.
Minor amounts of oxides other than Al2 O3 may also form in the BC/TC
interface [61]. Such oxides may, for example, form as a chromium rich layer
25
a)
b)
TC
Al2 O3
BC
Al2 O3
TC
2 μm
BC
1 μm
Figure 18: Protective layers of thermally grown Al2 O3 in the BC/TC interface.
a) A torn off top coat reveals the interfacial TGO. b) A cross-section showing a
layer of Al2 O3 .
between the Al2 O3 and TC or as bulky clusters containing a mixture of several types of oxides: (Al, Cr)2 O3 , Ni(Al, Cr)2 O4 and NiO [61]. Such clusters of
chromia–spinel–nickel oxide may form quite early during oxidation, and form
in greater quantities with higher temperature, but may remain fairly constant
once formed [61].
During oxidation, the oxide can form either internally, as a subscale, or as
an external scale, explained in Fig. 19. In order for the oxide layer to be protective, it must be external; hence, in a Ni–Cr–Al system, the formation of an external Al2 O3 scale must be promoted. There are several factors that influence the
ability of the alloy to form external alumina: oxygen partial pressure, amount
of solved O in the alloy, amount of aluminium in the alloy and the amount of
other alloying elements, most importantly Cr.
The effect of Al content and O concentration in the alloy surface can be understood by the following equation which gives the thickness, x, of the subscale
at time t in s [11].
2NO D O t
x=
νN M
µ
¶1
2
(4)
NO is the mole fraction of oxygen in the metal at the surface, D O is the diffusivity
of oxygen in the alloy, ν is the ratio of oxygen to metal atoms of the formed
oxide and N M is the mole fraction of the oxide forming element (Al in the Ni–Al
system). Fig. 20 shows the internal oxidation depth at 1000 °C as function of Al
content for different O concentrations at the alloy surface: max solubility of O
in Ni, and 50 %, 20 % and 5 % of full solubility [62]. As can be seen, the subscale
thickness decreases with increasing Al and decreasing O in the alloy; eventually,
26
a)
b)
atmosphere
atmosphere
oxide
x
alloy
oxide
alloy
Figure 19: Two types of oxidation: a) Internal oxidation: formation of a nonprotective subscale. b) External oxidation: formation of a protective oxide
scale.
Internal oxidation depth, µm
15
max solubility of O
50 % of max O solubility
20 % of max O solubility
5 % of max O solubility
internal to external
10
5
0
0
5
10
15
20
25
30
35
40
Al content, at.%
Figure 20: Subscale thickness as function of Al and O content. The hypothetical
transition from internal to external oxidation is also marked.
a shift to external oxidation will occur. For Ni–Al, the amount of Al needed to
cause a shift from internal to external oxidation is & 17 wt.% [10] as evident
from Fig. 12. The transition from internal to external oxidation may occur for
[51]
s
NM >
πg NO D O Val l oy
2νD M Voxi d e
(5)
where g is the fraction of formed oxide at which the internal subscale becomes
continuous and rate controlling, D M is the diffusivity of the oxidising element
27
Figure 21: The Al activity in a Ni–Cr–5 wt.% Al system as function of Cr content.
Data from Thermo-Calc.
and Val l oy and Voxi d e are the molar volumes of the alloy and the oxide respectively. Fig. 20 shows hypothetical transitions from internal to external scales
with g arbitrary set to g = 0.3 and D M taken as a rule-of-mixture mean of the
diffusivity in γ with 30 % β.
The addition of Cr to the Ni–Al system may also promote the formation of
external Al2 O3 through several mechanisms. As shown in Fig. 12, the addition
of 5 wt.% Cr, (∼ 5 at.%) enables the alloy to form Al2 O3 at an Al content as low
as ∼ 5 wt.%, (∼ 10 at.%). Chromium may, for example, act as a getter for oxygen
[63] which lowers the O concentration at the alloy surface. As evident from
Eq. 5 and Fig. 20, lowering NO makes it possible to form external Al2 O3 at lower
Al contents. Another effect of Cr addition is its influence on Al activity. This
is illustrated in Fig. 21 for a Ni–Cr–Al system with 5 wt.% Al at 1000 °C. The Al
activity has been calculated by Thermo-Calc as function of Cr content; it can
be seen that the addition of Cr increases Al activity.
3.2 The reactive element effect
Bond coat alloys contain minor additions of reactive elements, such as Y, Hf, Zr,
Ce or La [13, 22, 30, 36]. REs are generally considered to improve the oxide scale
adhesion; several mechanisms have been suggested:
• REs tie up sulphur which would otherwise have segregated to the metal/oxide interface and lowered the metal/oxide adhesion [30]. Lowering the
S content in the alloy can have the same effect [64].
28
• REs may slow down oxidation by segregating to Al2 O3 grain boundaries
and slow down Al grain boundary diffusion [65]. REs thus alter the oxide growth mechanism from an outward growing to an inward growing
oxide scale [35]. This also decreases spalling of the oxide by decreasing
lateral growth of the oxide, which could have happened if simultaneously
inward diffusion of O and outward diffusion of Al had occurred [35].
• REs may form oxides in the metal/oxide interface and mechanically pin
the oxide to the metal by so called pegging [49].
Y, which is the most common RE, readily forms oxides and may be found in
the Al2 O3 scale as: yttria Y2 O3 , yttrium aluminium perovskite (YAP) YAlO3 , and
yttrium aluminium garnet (YAG) Y3 Al5 O12 [49].
3.3 Breakdown of the protective oxide scale
The TGO will remain protective only as long as the bond coat contains enough
Al to maintain a continuous alumina scale. During high-temperature exposure, aluminium will be depleted through oxidation and interdiffusion with the
substrate [12, 37, 58]. An aluminium concentration of Ê3–5 wt.% is generally
enough to maintain the Al2 O3 scale [10, 11, 66, 67]; for low Al contents, nonprotective oxides may start to form in the BC/TC interface and the oxidation
rate increases; this marks the onset of breakaway oxidation, or chemical failure.
The chemical failure can be divided into two types: mechanically induced
chemical failure (MICF) and intrinsic chemical failure (InCF) [68]. MICF typically occurs during thermal cycling where the protective oxide scale cracks on
cooling and needs to be reformed; failure occurs when the Al content is too low
to heal/reform the protective alumina layer.
InCF occurs when the Al content beneath the oxide layer drops to such a low
level that the Al2 O3 is no longer the thermodynamically preferred oxide. This
occurs at considerably lower Al contents then MICF. This results in the formation of other oxides, either from the alloy or by decomposition of the alumina
scale according to reactions such as [58]:
Al2 O3 + 2 Cr
Cr2 O3 + 2 Al
(F)
or
Al2 O3 + 21 O2 + Ni
NiAl2 O4
(G)
The Al2 O3 scale is thus replaced, or partially replaced, by a layer of chromia
(Cr, Al)2 O3 , spinel (Ni, Co)(Cr, Al)2 O4 , nickel oxide and cobalt oxide [40, 50, 53,
29
69–71]. Internal oxidation of the remaining aluminium may also occur [69].
These TGOs are not as protective as alumina and the layer of chromia and
spinels has lower interfacial fracture resistance which may cause the top coat
to spall on cooling [12, 40, 68, 69].
30
4
Thermal fatigue of coatings
In addition to applied mechanical load, there are two sources for stresses in an
APS TBC system: 1) growth stresses in the interface TGO and 2) stresses that
develop on heating or cooling due to the mismatch in coefficient of thermal expansion between the bond coat, interface TGO and top coat [30]. Both sources
of stress act at the bond coat/top coat interface and failure of TBC systems consequently occurs by fracture in, or close to, the BC/TC interface [12].
Oxide growth stresses can partly be understood from the so called PillingBedworth ratio, P B R, which is calculated as [72]
Wd
(6)
wD
where d and D are the densities of the metal and the oxide respectively, and w
is the amount (weight) of metal necessary to produce the amount (weight) W
of oxide. A P B R < 1 gives tensile stresses in the oxide while a P B R > 1 gives
compressive stresses (more so the higher the PBR). Table 2 shows the P B R for
some common oxides in coatings [51].
PBR =
Table 2: P B R of some oxides
common in coatings [51].
oxide
PBR
Al2 O3
Cr2 O3
NiO
CoO
1.28
2.07
1.65
1.86
Gas turbine starts and stops give cyclic variations in temperature and the
resulting cyclic thermal stresses make the TBC system susceptible to fatigue.
31
The thermal mismatch stresses are often considered to be most harmful during
cooling [70, 73] as, during heating, stress relaxation may occur. During cooling, however, there is little time for stress relaxation and stresses develop in the
BC/TC interface that depend on the temperature drop [70, 73]; for large temperature drops, the thermal mismatch stresses during cooling may dominate
over the TGO growth stresses [30].
The stresses that develop due to thermal mismatch depend, in addition to
the temperature drop and the CTE mismatch, on BC/TC interface morphology
and the thickness of the interface TGOs [74]. The TGO thickness and composition change as the TBC system is exposed to high temperature and thus affect
the BC/TC interface stresses.
A simplified description of the stresses that arise during thermal cycling of
a TBC system would be as follows [75]: When the TBC system is heated to high
temperature, stresses are introduced in the BC/TC interface due to the differences in CTE between the bond coat and the top coat. These stresses are, partly
or entirely, reduced due to stress relaxation at high temperature [73]. Long
high-temperature exposure causes the interface TGO to grow, resulting in TGO
growth stresses, which may also relax at high temperature. At cooling, stresses
are again introduced due to differences in the CTE, only now there is little time
for stress relaxation and stresses develop in the BC/TC interface.
In a rough BC/TC interface, with alternating peaks and valleys, thermal cycling would cause out-of-plane tensile stresses to form at the interface peaks
and out-of-plane compressive stresses at interface valleys, as shown in Fig. 22 a).
As the interface TGOs grow, the stress distribution will be affected as illustrated
by Fig. 22 b) and c). A thicker layer of interface TGOs causes the compressive
stresses at the valleys to shift to tensile stresses. Stresses formed through this
mechanism will be able to propagate fatigue cracks in the vicinity of the BC/TC
interface, and, consequently, cause the TBC system to fail by fatigue.
a)
b)
-
+
BC
as-sprayed
-
c)
TC
TC
-
+
BC
+
4 μm TGO
TC
-
+
BC
+
8 μm TGO
Figure 22: Out-of-plane (vertical) stresses in the BC/TC interface. The compressive stresses at the valleys shift to tensile stresses as the TGO grow. Based
on Jinnestrand and Sjöström [75]
32
4.1 Crack nucleation mechanisms
The plasma spray process gives the top coat a very characteristic splat-on-splat
structure. The degree of inter-splat bonding can be rather modest which gives
rise to many crack-like defects in the top coat, see Fig 15. These pre-existing
interlamellar delaminations in the top coat may act as crack embryos. Several
papers have attributed crack nucleation to these pre-existing interlamellar delaminations [50, 76–81].
Cracks have also been described to nucleate in the interfacial TGO during
cycling. Crack initiation in the BC/TC interface is most commonly attributed
to peak and off-peak positions in the BC/TC interface [48, 82, 83]. There are
several suggested explanations for this, such as:
1. Out-of-plane tensile stresses prevail at peak and off-peak positions.
2. During cycling, in-plane compressive growth stresses may cause the layer
of interfacial Al2 O3 to buckle and delaminate at peak positions. The TGO
will reform and the process is repeated. Such repeated delamination and
regrowth will give rise to a layered TGO structure at peak positions, shown
in Fig. 23 a), which may act as starting points for larger delamination
cracks [36, 82, 84, 85].
3. Cracks may nucleate in the TGO due to large growth stresses in the voluminous clusters of chromia and spinels, shown in Fig. 23 b), that may
form in addition to Al2 O3 during high temperature exposure [50, 61, 80,
81].
a)
b)
TC
TC
TGO
BC
20 μm
BC
TGO
10 μm
Figure 23: Crack formation in the interfacial TGO. a) Repeated cracking and
regrowth giving a layered structure in the TGO. b) Cracking in a cluster of chromia, spinels and nickel oxide.
33
4.2 Crack growth mechanisms
There are several micromechanic failure mechanisms described in literature;
they differ in where the crack is assumed to have nucleated and the path of
the growing crack, Fig. 24, 25 and 26 summarise some of these. Many of the
suggested failure mechanisms from the literature assume that crack nucleation
and growth mechanisms act on a microscale, typically the scale of an interface
peak and valley. It is then assumed that crack growth, as that on a microscale,
occurs simultaneously throughout the BC/TC interface such that failure eventually occurs by coalescence of microcracks causing the top coat to spall [12].
It hence becomes sufficient to study crack growth only in a least representative
cell. The BC/TC interface is often idealised as some periodic function, often a
sine wave, and the least representative cell typically includes a peak and a valley, similar to those shown in Fig. 24, 25 and 26.
The failure mode shown in Fig. 24 assumes crack initiation at peaks in a
sinusoidal BC/TC interface, Fig. 24 a). The crack path will then either follow the
BC/TC interface, Fig. 24 b), or kink out in the TC, Fig. 24 c) [73, 86]. Such crack
growth is assumed to occur at every peak in the BC/TC interface and failure
occurs when such microcracks meet and coalesce.
Fig. 25 shows a failure mode where crack nucleation occurs by opening of
the pre-existing interlamellar delaminations in the top coat. Such crack embryos grow in the vicinity of BC/TC interface peaks and arrest as they encounter
a BC peak, Fig. 25 a). Meanwhile, the growth of the interface TGO will increase
the out-of-plane tensile stresses at off-peak positions, and when such stresses
are high enough, cracks will nucleate in the TGO at peak positions, Fig. 25 b).
The crack propagation then proceeds until several of these cracks coalesce and
cause failure, Fig. 25 c) [76, 77].
Another mechanism, similarly, suggests that cracks initiate from pre-existing delaminations in the top coat, but above the bond coat peak positions,
Fig. 26 a). Such cracks initiate while the out-of-plane stresses are tensile at peak
positions but still compressive at valley positions, see Fig. 22 a). Since the cracks
cannot grow through the areas of compressive stresses at flank and valley positions, the crack arrests until the growth of the TGO changes the compressive
flank and valley stresses to tensile stresses as illustrated by Fig. 22 b) and c).
Crack growth occurs in the top coat and the failure occurs as these cracks coalesce, Fig. 26 b) [78, 79].
In addition to failure by the growth and coalescence of microcracks, failure
may also occur through growth of macrocracks from the edges of TBC coated
specimens [87, 88]. Edge cracking occurs due to the stress concentration at
the TBC edge where the coating ends. Sjöström and Brodin [87] investigated
the influence of the chamfer angle on the edge cracking of TBCs and found
34
a)
crack
TGO
b)
top coat
TGO
bond coat
c)
top coat
crack
crack
TGO
bond coat
top coat
bond coat
Figure 24: Crack nucleation in the TGO, a), followed by either: b) crack growth
in, or close to, the TGO, or c) crack growth by kinking out in the TC.
a)
crack
TGO
b)
top coat
TGO
bond coat
c)
top coat
crack
crack
TGO
bond coat
top coat
bond coat
Figure 25: Crack nucleation in the top coat, a), followed by b) damage of TGO
and c) crack growth.
a)
crack
TGO
top coat
bond coat
b)
top coat
crack
TGO
bond coat
Figure 26: Crack growth in the top coat: a) nucleation and b) growth.
o
60
top coat
edge cracking
o
90
top coat
chamfer angle
substrate
0
bond coat
Figure 27: Edge cracking of TBCs.
35
that any chamfer angle larger than 60◦ gave essentially the same risk of edge
cracking whereas an angle < 60◦ gave a lower risk of edge cracking, see Fig. 27.
The crack growth paths suggested in literature is evidently quite divers, ranging from entirely in the BC/TC interface to mixed interface/TC fracture to entirely in the TC. Fracture that occurs in the TC is referred to as white fracture
and fracture that occurs in the BC/TC interface is called black fracture as the
fracture surfaces will appear white and black respectively when examined optically. Fracture by cracks that grow partly in the BC/TC interface and partly in
the TC is referred to as mixed fracture.
Which failure mode, white, black or mixed, that results depends partly on
the thermal cycle. Trunova et al. [89] compared the fracture mode of TBC systems subjected to thermal cycling with different dwell times. It was found that
long dwell time gave mainly black fracture, intermediate dwell times gave mixed
fracture and no dwell time gave white fracture. It was also noted that, as the
dwell time decreases, the total number of cycles to failure increased and the
cumulative high temperature exposure time decreased.
4.3 Coating life assessment
4.3.1 Microstructure-based life models
There are several microstructure- or composition-based life models for TBCs;
they all rely on the tendency of TBC life to depend on TGO growth. Rabiei and
Evans [77] showed, for example, that a TGO thickness in excess of ∼ 5.5 µm
could be critical for TBCs. TBC life can thus be predicted from oxide growth
modelling.
Other models are in different way based on aluminium depletion of the
bond coat. As the aluminium content reaches a critical value, the coating stop
acting as an Al reservoir; the protective interface oxide layer can no longer be
maintained and chemical failure occurs [12, 39, 49, 66, 90–93]. By modelling
oxidation and interdiffusion with the substrate, TBC life models have been established [66, 90, 92–99].
It is necessary to have a criterion for when the aluminium content becomes
too low. Since removal of Al from the coating causes dissolution of the β phase,
the complete depletion of β phase, which can be observed from micrographs of
cross-sectioned specimens, can be used as a life criterion [12, 91–93]. In other
cases, a critical Al content may be a more suitable criterion; Renusch et al. [66]
suggested ∼ 3 wt.% as a critical Al content.
36
4.3.2 The NASA model
The NASA model [100] takes its starting point in a Coffin-Manson type expression
µ
N=
∆εi
∆ε f
¶b
(7)
where N is cycles to failure, ∆εi is the inelastic strain range, ∆ε f is the inelastic
strain range that causes failure in one cycle and b is a constant. The effect of
high temperature exposure on life is included in ∆ε f which is a function of TGO
growth and the adhesion strength of the unoxidised TBC. The growth of the
TGO influences life through the expression
µ
¶
µ ¶d
δ c
δ
∆ε f = ∆ε f 0 1 −
+ ∆εi
δc
δc
(8)
where ∆ε f 0 is the inelastic failure strain range for an unoxidised coating system,
δc is the critical oxide layer for which the coating would fail in a single cycle
and c and d are constants which are c ≈ d ≈ 1. The oxide thickness, δ, can be
obtained from a power-law equation, see equation 3.
4.3.3 Model suggested by Busso et al.
Busso et al. [76] have suggested the following life model for APS TBCs:
dD = D m
³σ
´
max p
F
dN
(9)
where 0 É D É 1 is a fatigue damage parameter such that D = 1 at failure, σmax
is the maximum out-of-plane interfacial stress, N is number of cycles and m
and F are given by
σmax
m = 1 −C
σc0
µ
¶0.818p
(10)
and
F = F 0 (1 − F 1 σmax )
(11)
37
where σc0 is the initial strength of the TBC and p, C , F 0 and F 1 are material
parameters that need to be calibrated to experimental data.
The fatigue damage is driven by the maximum out-of-plane stress, σmax ,
which is obtained from finite element analysis of a least representative cell of a
TBC system. The σmax includes the combined stresses from the thermal cycle,
oxide growth and sintering of the top coat according to [76, 101]
σmax = σtherm. + σox. + σsintr.
(12)
where σtherm. , σox. and σsintr. are functions of temperature cycle, cumulative oxidation time and BC/TC interface morphology; they describe the out-of-plane
stress contributions from thermo-elastic and visco-plastic deformation, oxidation and sintering.
4.3.4 Model suggested by Brodin, Jinnestrand and Sjöström
Brodin, Jinnestrand and Sjöström [70, 73, 102] have developed a fracture mechanics based life model using a Paris law type of equation:
dD
= C (λ∆G)n
dN
(13)
where G is the energy release rate and C and n are constants. D is a damage
parameter which can contain contributions from cracks in the BC/TC interface,
, and cracks running partly in the top coat and in the
l iTGO , in the top coat, l TC
j
BC/TC interface, l kTC/TGO , according to
P
D=
TGO P TC P TC/TGO
+ j l j + k lk
i li
L
(14)
where L is the total analysed length.
The model involves cracks that partially, or completely, follow the BC/TC
interface; such cracks will grow in a mixed mode [73]. To account for mixed
mode crack growth, G is multiplied by a mixed mode function, λ [103]:
λ = 1 − (1 − λ0 )
µ
µ
¶¶
2
∆K I I m
tan−1
π
∆K I
(15)
where ∆K I and ∆K I I are the stress intensity factors in mode I and II, and λ0 and
m are constants.
38
The influence of thermal loads, surface morphology and interface TGO growth is included in the finite element calculation of ∆G and ∆K II /∆K I [73]. ∆G
is calculated by a virtual crack extension method. ∆K II /∆K I is calculated from
the crack flank displacements using the theory of interface cracks suggested by
Hutchinson and Suo [103].
39
5
Experimental methods
5.1 Thermal fatigue
When a TBC system is cycled, the difference in coefficient of thermal expansion between the bond coat and the top coat causes thermal stresses close to
the BC/TC interface. In the lab environment, two main types of thermal cycling tests exist: thermal cycling fatigue (TCF) (or furnace cycle test (FCT)) and
burner rig test (BRT) (or thermal shock). Brodin [86] found, from literature,
that TCF tests often had average heating and cooling rates in the order of 1–
2 °C, whereas BRT tests often had average heating and cooling rates in the order
of 6–12 °C. BRT does also, to a greater extent, give white fracture whereas TCF
gives black fracture [86].
The burner rig test uses a flame to heat the specimen on the coated side.
Burner rigs typically reach a maximum gas temperature of 1350–1750 °C [104].
During heating, the specimens are sometimes cooled on the uncoated side to
introduce a larger temperature gradient during heating. After heating, the specimens are rapidly cooled by compressed air, typically with no, or very short,
hold time at high temperature. Fig. 28 shows a schematic drawing of a burner
rig and Fig. 29 a) gives an example of a BRT temperature curve. In addition to
thermal shock testing, burner rigs can also be used for oxidation tests with long
dwell times and hot corrosion tests, the latter typically performed at temperatures around 900 °C [104].
In a furnace cycle test, the specimens are thermally cycled by moving in
and out of a furnace. Such testing is associated with lower heating rates than
the burner rig test and the temperature gradients in the specimens are lower
[86]. Furthermore, the high-temperature dwell time is usually very long compared to BRT. During cooling, the specimens can be cooled by compressed air.
41
Fig. 30 shows a schematic drawing of furnace cycling and Fig. 29 b) shows a
temperature curve.
a)
b)
TBC
TBC
airflow
airflow
flame
fixture
fixture
Figure 28: Schematic drawing of a burner rig. a) Heating by a flame on the
coated side while cooling the uncoated side. b) The specimen is moved out of
the flame and cooled by air.
a) 1100
heating
900
cooling
800
b) 1200
1000
temperature, ◦ C
temperature, ◦ C
1000
top coat temp.
substrate temp.
700
600
cooling
600
400
200
500
400
heating
800
0
20
40
60
80 100
time, s
120
140
0
160
0
10
20
30
40
time, min
50
60
70
Figure 29: Two thermal cycles. a) A burner rig cycle, based on Liu et al. [88]. b)
A furnace cycle with forced air cooling.
a)
b)
furnace
airflow
specimens
Figure 30: Schematic drawing of a cyclic furnace. a) Dwelling in furnace during
the hot part of the cycle. b) Cooling with air during the cold part of the cycle.
42
5.2 Corrosion test
Corrosion resistance evaluation of coatings may be conducted in many ways.
Often, salt mixtures containing Na2 SO4 are deposited onto the specimens in
amounts of 0.075–5 mg/cm2 and in intervals of 20–100 h. Salt deposition can
occur with or without the addition of gaseous SO2 in the temperature interval
700–1000 °C [105–111]. Table 3 summarises some of the corrosion conditions
encountered in literature for coatings [105–107, 110–112]. Water vapour has
also shown to influence the corrosion rate. Testing conditions including 5–15 %
water vapour at 800–1100 °C can be found in literature [64, 113–117].
A cyclic corrosion rig was built during the duration of the project to evaluate the corrosion properties of coatings. Fig. 31 shows the cyclic corrosion rig
which operates by moving the furnace up and down. When in the lower position, the furnace covers the specimens; when in the upper position, the specimens are allowed to cool slowly in air; no forced airflow was used for cooling
moving lid
specimens
nozzles
moving furnace
furnace
water inlet
Figure 31: Cyclic corrosion rig.
43
Table 3: Some test parameters for corrosion, from literature [105–107,
110–112].
Salt
Amount salt,
mg/cm2
Interval of salt
deposition, h
Total test
duration, h
0.075–0.75
2.5
3–5
2
1
0.1
50
24
51
20
100
20
500
168–312
51
100
300–500
100–430
Na2 SO4 –20 mol% K2 SO4
Na2 SO4 –0–5 wt.% NaCl
Na2 SO4 –50 wt.% V2 O5
Na2 SO4 –25 wt.% NaCl
Na2 SO4 –K2 SO4
Na2 SO4 –10 wt.% NaCl
Temp.,
°C
700–900
900
800
900
950
1000
800
700
°
Temperature, C
600
500
400
300
200
100
0
47.6
47.8
48
48.2
48.4
Time, h
48.6
48.8
49
49.2
Figure 32: A temperature cycle during cyclic corrosion; the sudden drop in temperature at the end of the cooling is due to salt-solution deposition.
in this rig. When the furnace is in its upper position, a lid closes to reduce heat
loss in the furnace; attached to the lid are also a number of nozzles which sprays
the specimens with a corrosive substance. During the hot part of the cycle, the
option exists to spray water into the furnace which instantaneously evaporates.
Fig. 32 shows the temperature curve measured on a specimen during a corrosion cycle.
5.3 Adhesion test
A commonly used adhesion test procedure is to attach the coated and uncoated
sides of a button specimen to two bars, that can be mounted in a tensile test
machine, and load the coating until fracture. The set-up is schematically shown
in Fig. 33. The tensile test machine must be equipped with universal joints to
ensure moment free mounting. The specimen is fastened to the bars by a suitable adhesive, most commonly by epoxy which is cured at moderate temper44
atures (120–175 °C [19]). During curing, a slight compressive load is applied
to the bar/specimen/bar system to ensure good adhesion between fixture and
specimen. The method is described in ASTM C633 Standard test method for
adhesion or cohesion strength of thermal spray coatings [118] as well as EN 582
Thermal spraying – determination of tensile adhesive strength [119].
bar
adhesive
coating
substrate
adhesive
bar
Figure 33: Set-up for adhesion testing of TBCs. Based on Davis [19].
To ensure good results, the specimens need to be flat and the surfaces need
to be clean and free from loose material; the coating may therefore have to be
ground or grit blasted. Furthermore, any coating overspray onto the sides of
the button specimen, as well as beads of excessive adhesive at the joint, must
be removed before testing. Even if great care is taken, some critical comments
to the method are [19]:
• Bending moments introduced by mounting in the tensile test machine
will give erroneous results. This is avoided by the use of a self-aligning fixture, as shown in Fig. 33, and by ensuring that the specimens are ground
flat before bonding to the fixture.
• The type of adhesive will influence the result, as will the thickness of the
adhesive film. In a porous coating, the penetration depth of the adhesive
will affect the results. These effects stress the importance of a consistent
and repeatable curing procedure.
45
• The strength of the adhesive sets the upper limit for how strong coatings
can be tested.
• The tensile test is unsuitable for evaluation of very thin and very porous
coatings.
• Variation in coating thickness, distribution and size of defects in the coating and residual stresses may give scattered data.
5.4 Microscopy
5.4.1 Specimen preparation
TBC systems are challenging to prepare for microscopy since the ceramic top
coat and TGOs may crack during polishing. Care must be taken not damage the
ceramic parts during specimen preparation since [19]:
1. Ceramic particles that come loose during polishing, pull-outs, may scratch the specimen.
2. Pull-outs in the top coat will make the top coat appear more porous, so
that the porosity of the top coat will be overestimated.
3. Cracking of the top coat and TGO during polishing makes it impossible
to determine the amount of damage that was caused by thermal cycling
of the TBC system.
To avoid damaging the TBC during specimen preparation, the TBC system
is infiltrated with epoxy in vacuum. The specimen is placed in vacuum and
epoxy is poured over the specimen. The epoxy will penetrate the pores and
cracks of the top coat and reduce the amount of damage introduced during
the subsequent cutting, mounting and polishing. Would the top coat be damaged during the polishing, it is still possible to distinguish damage introduced
by the thermal cycling as those cracks would be filled with epoxy and cracks
introduced during polishing would not.
5.4.2 Scanning electron microscopy
Scanning electron microscopy (SEM) offers, in addition to imaging, micro-analysis tools such as energy dispersive spectroscopy (EDS) and wavelength dispersive spectroscopy (WDS), as well as electron backscatter diffraction (EBSD),
a tool for studying crystallographic structure and orientation.
46
EDS and WDS provide compositional information of the specimen. WDS
can better resolve different elements in the collected X-ray spectrum, but is
more time consuming than EDS. The spatial resolution of both techniques depends on the electron beam energy and the studied element according to [120]
RK O =
0.0276A 1.67
E
Z 0.89 ρ 0
(16)
where R K O is the electron range (the X-ray generating volume) in µm, A is the
atomic weight in g/mole, Z is the atomic number, ρ is the density in g/cm3 and
E 0 is the electron beam energy in keV. With E 0 = 20 keV, EDS/WDS on Ni would
give a spatial resolution of ∼ 1.4 µm and Al would give ∼ 4.2 µm. Poor spatial
resolution of these techniques may sometimes be problematic.
A problem with SEM on TBC systems is that the top coat and the TGOs
do not conduct electricity which causes the specimen to charge. This can be
solved by applying a thin carbon film onto the surface of the polished specimen.
5.5 Interface roughness measurement
The BC/TC interface roughness can be measured on cross-sectioned specimens by image analysis. A Matlab script was written for the purpose. The interface roughness profile is acquired from greyscale light-optic micrographs. The
steps of the acquisition process are outlined in Fig. 34 and involves:
1. Thresholding of the grayscale image to obtain a binary image, see Fig. 34
a) and b).
2. Acquisition of the interface roughness profile by simulating a profilometer stylus tip, see Fig. 34 c).
3. Filtering of the profile and establishing the mean line, see Fig. 34 d).
4. Cropping the profile to a suitable sampling length, see Fig. 34 e).
The interface roughness profile is then used for calculation of various surface roughness parameters, the most well-known probably being the profile
arithmetic mean deviation, Ra:
Ra =
1
l
Z
0
l
|z(x)|dx
(17)
47
where l is the analysed length and z and x are explained by Fig. 34 e). A comparison of Ra values, for some different surfaces, showed that the Ra values
obtained by image analysis were in good agreement with those obtained by a
profilometer.
a)
b)
105
z, μm
100
90
z, μm
c)
85
160
140
120
100
80
0
100
200
300
400
200
300
400
acquired profile
stylus tip
95
metal/ceramic interface
425
430
435
440
x, μm
445
500
600
700
800
900
500
600
700
800
900
500
600
700
800
x, μm
450
d)
z, μm
50
0
−50
0
profile
mean line
100
x, μm
e)
z, μm
50
0
−50
100
200
300
400
x, μm
Figure 34: Roughness profile acquisition procedure: a) greyscale image, b) binary image, c) stylus tip simulation, d) filtering and establishing mean line, and
e) the obtained roughness profile.
48
6
Discussion of appended papers
Paper I: Influence of isothermal and cyclic heat treatments on the adhesion of plasma sprayed thermal barrier coatings
The influence of different heat treatments on the adhesion properties of APS
TBCs was studied with the aim of relating the adhesion strength of the TBC to
high-temperature exposure and to the test method used. The studied TBC system consisted of a Hastelloy X substrate coated with an APS NiCoCrAlY bond
coat and APS Y-PSZ as top coat. The specimens were subjected to three different thermal treatments: isothermal oxidation, furnace cycling and burner rig
test; all at a temperature of ∼ 1100 °C. The adhesion of the coatings after thermal treatment was evaluated by the ASTM C633 method.
The differences between the thermal exposure conditions affected the microstructure. Specimens subjected to the burner rig test developed a very thin
TGO and were never depleted of the Al-rich β-phase. Furnace cycling and
isothermal oxidation developed thicker TGO and were depleted of β-phase after ∼ 100 h. Furnace cycling and isothermal oxidation also gave more pronounced interdiffusion with the substrate.
The adhesion of the isothermally heat treated specimens increased about
50 % compared to specimens tested in the as-sprayed condition. Isothermally
oxidised specimens were also found to retain their adhesion strength with oxidation time whereas the cyclic tests both decreased the adhesion with number of cycles. The adhesion decrease of the two cyclic heat treatments was attributed to the fatigue damage introduced in the specimens during cycling.
The paper provides valuable information on how different test methods influence the degradation of TBCs. If the decrease in adhesion is taken as an
49
indication of damage introduced in the TBC, then the tests, in increasing order
of severity, can be ranked as 1) isothermal oxidation, 2) furnace cycling and 3)
burner rig test, for the same exposure time. However, the burner rig test gave a
slower decrease of adhesion per cycle: ∼ 1.54 · 10−3 MPa/cycle, compared with
∼ 20.44 · 10−3 MPa/cycle for furnace cycling. Hence, if the life is measured in
number of cycles, as opposed to time, the burner rig test gives longer lives than
furnace cycling. This is consistent with the tendency for longer high temperature dwell times to reduce the number of cycles to failure.
Paper II: Fractographic and microstructural study of isothermally and cyclically heat treated thermal barrier
coatings
A fractographic study was performed on fracture surfaces from adhesion tested
TBC systems exposed to different heat treatments. The studied TBC system
consisted of a Hastelloy X substrate coated with APS NiCoCrAlY and APS YPSZ. The coated specimens were heat treated at a temperature of ∼ 1100 °C by
isothermal oxidation, furnace cycling and burner rig test. The top coats were
pulled off the specimens in a tensile test machine and the resulting fracture
surfaces were studied.
The fracture occurred either almost entirely in the top coat or both in the
top coat and the BC/TC interface (mixed fracture); no specimen fractured entirely in the BC/TC interface. Isothermal oxidation gave fracture almost entirely
in the top coat while the two cyclic heat treatments gave mixed fracture. However, even for cyclic heat treatment, > 80 % of the fracture occurred in the top
coat. The amount of BC/TC interface fracture did increase with number of cycles.
Regions where the fracture occurred in the top coat showed essentially the
same characteristics regardless of heat treatment; the fracture occurred mainly
between the splats in the splat-on-splat structure typical for APS coatings. Through-splat fracture occurred sparingly and was often associated with discontinuities in the microstructure such as partially melted particles.
If the amount of interface fracture from the adhesion test is taken as an indication of the amount of interface damage in the specimen, isothermally oxidised specimens were essentially undamaged at ∼ 300 h of exposure whereas
the furnace cycling test had developed some interface damage. The furnace cycle test had a larger increase in interface damage per cycle; it takes the burner
rig test ∼ 1000 cycles to reach the same damage as the furnace cycle test reaches
in ∼ 300 cycles. Both cyclic tests reached the same amount of damage roughly
50
at the same fraction of their total lives; from experience, ∼ 300 cycles and ∼
1000 cycles for TCF and BRT respectively corresponds to 30–50 % of the total
life. A microstructural study of the top coat also show that the short hold times
associated with burner rig tests gave considerable less sintering of the top coat.
The study illuminates several of the degrading mechanisms of TBCs during
isothermal and cyclic high-temperature exposure: fatigue damage, interface
TGO growth and sintering of the top coat. In particular, the study showed the
difference between isothermal and cyclic heat treatment. While the isothermal
oxidation gave fracture entirely in the top coat, cyclic heat treatments gave increasing fractions of BC/TC fracture with number of cycles.
Paper III: Influence of substrate material on the life of
atmospheric plasma sprayed thermal barrier coatings
The substrate material on which the coating is deposited may influence the life
of the TBC. Therefore, a study on the influence of substrate material on the life
of APS TBCs was conducted. The specimens consisted of Hastelloy X (HX) and
Haynes 230 (H230) substrates with an APS NiCoCrAlYSiTa bond coat and an
APS zirconia top coat. The specimens were exposed to both isothermal oxidation and thermal cycling at 1100 °C. The specimens were tested until failure and
the resulting fracture surfaces were studied. Cross-sections were also prepared
and used for measurements of the interdiffusion between substrate and bond
coat by EDS, as well as to measure the Young’s moduli of the bond coat and the
substrate by nanoindentation.
The H230 specimen had a thermal cycling life almost twice as long as the HX
specimen, 1070 cycles compared to ∼ 600 cycles. Investigation of the fracture
surface revealed that both the HX and H230 specimens had failed roughly at
the same fraction of black fracture, ∼ 40 %. The black fracture was also shown
to have occurred mainly in the TGO/TC interface rather than within the TGO.
During the isothermal oxidation test, the specimens were removed from the
furnace at regular intervals and the test was interrupted when the TBC spalled
during removal from the furnace. The HX specimen, again, had a shorter life,
1000 h compared to 1650 h for H230.
EDS line scans on cross-sectioned specimens provided through-coating compositional profiles; it was found that Fe, Mn, Mo and W from the substrate
diffused into the bond coat; the interdiffusion was more pronounced for the
HX specimen. EDS on the fracture surfaces revealed that Mn and Fe from the
substrate could be found in the oxide; Fe was present in very small amounts,
whereas Mn appeared to easily form oxides in the BC/TC interface. EDS results
51
also revealed that the H230 specimen retained an Al reservoir in the bond coat
centre longer than the HX specimen.
TGO thickness measurements revealed that, while the growth rates of Al2 O3
were similar for both the HX and H230 specimens, there existed a difference in
the total TGO growth rate (the total TGO growth rate also included oxides such
as Cr2 O3 , spinels and NiO). The HX specimen had a faster total TGO growth
than the H230 specimen. At the time of failure, the total TGO thickness was
rather similar for the two specimens: 16.5 µm and 14.4 µm for the HX and H230
specimens respectively.
The Young’s moduli, measured by nanoindentation, were used together with
CTE estimations, made by Thermo-Calc, in a finite element analysis of crack
growth in the BC/TC interface. It was found that the small differences in Young’s
modulus and CTE, which arose due to interdiffusion, did not cause any significant differences in the crack driving force and could thus not explain the difference in life. The difference in thermal cycling life of the two TBC systems
was considered to be due to differences in Al depletion and growth rate of nonAl2 O3 oxides. The faster total TGO growth rate for the HX specimen was likely
to have caused a shorter life due to higher stresses in the interface and the fact
that non-Al2 O3 oxides tend to crack and cause additional damage to the BC/TC
interface. The study showed that the substrate material influences the life of
the coating and that the growth kinetics of all oxides, not only Al2 O3 , may have
to be taken into account in life prediction of TBCs.
Paper IV: TBC bond coat-top coat interface roughness:
influence on fatigue life and modelling aspects
A large number of studies on TBC life prediction and TBC durability rely on finite element modelling of the stresses in the BC/TC interface. The influence
of BC/TC interface roughness on the thermal fatigue life of TBCs was therefore studied. Four TBC systems, with varying BC/TC interface roughness, were
thermally cycled to failure and their fatigue lives were correlated to their corresponding BC/TC interface roughness. Based on the results, a few different interface models for finite element analysis were developed. The specimens consisted of Haynes 230 coupons coated with VPS NiCrAlY and APS Y-PSZ which
were thermally cycled until failure.
The results showed that a rougher interface increased the fatigue life of APS
TBCs. It was found that a ∼ 50 % increase in Ra gave a ∼ 70 % increase in TCF
life. This increase in life was attributed to roughness, and not to oxidation effects, since no major differences in oxidation composition and kinetics were
52
observed.
A number of different roughness parameters were evaluated to establish
which parameters that best captured the characteristics of the interface. It was
found that an amplitude parameter together with a slope parameter accurately
captured the main features of the interface. The two parameters chosen as basis for interface model formulation were Rq and R∆q.
A number of interface models were created based on the selected interface
roughness parameters. The interface models were evaluated by finite element
analysis and by the life model suggested by Brodin, Jinnestrand and Sjöström,
(described in section 4.3). The life model initially predicted the smoothest interface to have the longest life which contradicted experimental observations.
It was concluded that the crack growth path must be adjusted for surface roughness effects; experimental observations suggested that higher roughness would
shift the crack growth path so that the cracks to a greater extent grew in the top
coat as opposed to in the BC/TC interface. When the finite element model was
adjusted to agree with the observation, the life model was able to capture the
experimentally observed results of high roughness promoting longer lives.
The study illustrated the large influence of BC/TC interface roughness on
the fatigue life of TBCs, and stressed the importance of accurate interface models in life prediction of TBCs. It was found that a sinusoidal interface model
based on an amplitude parameter and a slope parameter may be sufficient for
modelling of stresses in the interface.
Paper V: Cyclic hot corrosion of thermal barrier coatings and overlay coatings
Coatings in gas turbines may degrade due to corrosion. Testing was performed
with the aim of mapping the relative severity of: 1) oxidation in lab air, 2) oxidation in moist air, 3) corrosion with salt deposits in lab air, and 4) corrosion with
salt deposits in moist air. Both NiCoCrAlYSiTa coated and NiCoCrAlYSiTa + TBC
coated specimen were used. The testing was cyclic, including 48 h at 750 °C and
30 min cooling to a minimum of ∼ 100 °C. For the moist air condition, water was
injected in intervals of a few minutes during the hot part of the cycle; and for
salt deposition, a salt–water solution was sprayed on the specimens automatically at the end of the cold part of the cycle. Prior to analysis, the specimens
were cleaned to remove the residues from the salt deposition.
For NiCoCrAlYSiTa coated specimens, salt deposition gave notably thicker
TGOs. The TGOs of the salt exposed specimens also contained an Y-rich oxide
which was not observed for oxidation in lab air or moist air. EDS on the oxi53
dised/corroded surfaces revealed that the TGOs on all specimens consisted of
predominately Al2 O3 . Very low amounts of S were found on the surfaces indicating that there was little or no sulfidation; the presence of Na in the TGOs was
also low. The apparent influence of water was to change the TGO morphology.
For TBC coated specimens, the salt penetrated the top coat and sped up
oxidation/corrosion at the BC/TC interface. The salt exposed specimens developed a, at least locally, thicker TGO and the metal–oxide front of attack was
considerably more uneven than for oxidation in lab air and moist air.
The performed corrosion test gave an overview of the relative influence of
water vapour and salt deposits on the corrosion resistance of coatings. It also
showed that bond coat corrosion due to salt deposition may be present also for
TBC coated specimens. While the performed corrosion test may be unsuitable
for careful corrosion kinetics measurements, the present test parameters could
be used for screening and ranking of coatings.
Paper VI: Modeling of microstructural evolution and
lifetime prediction of MCrAlY coatings on nickel based
superalloys during high temperature oxidation
Aluminium depletion based life models involve oxidation and interdiffusion
modelling; thermodynamics software such as Thermo-Calc and DICTRA show
promising results when used for such purposes. Therefore, DICTRA was used
for the development of an Al depletion based life model for MCrAlY coatings.
The modelling was based on HVOF CoNiCrAlYSi coatings deposited on substrates of Inconel 792. The specimens were subjected to isothermal oxidation at
900–1100 °C and thermal cycling at 1100 °C. The interdiffusion simulation was
made in DICTRA and the oxidation simulation was made by a Matlab script
written for the purpose. The DICTRA script and Matlab script were alternately
run in an iterative process, the result of which was a coupled oxidation–interdiffusion model for prediction of Al depletion in the coating. The aim of the
study was to model the depletion of the β-phase, i.e. the decrease in thickness
of the β-containing zone in the middle of the coating. The end of the life of the
coating was taken to occur when the β-phase had been completely depleted.
The diffusion in a multi-phase system in DICTRA requires the calculation
of an effective diffusivity which depends on the geometrical phase distribution. Several geometrical phase distributions were tried, but a simple rule-ofmixture model was found appropriate. Included in the model was also a diffusion blocking effect. This occurred due to oxides, pores and grit residues at
splat boundaries and at the substrate/coating interface. The diffusion blocking
54
was introduced by correction of the diffusion time.
The coating developed two β-depleted zones, an inner, due to interdiffusion with the substrate, and an outer, due to oxidation. The inner and outer
β phase depletion occurred faster for higher oxidation temperatures. For the
outer zone, thermal cycling was found to give faster depletion than isothermal
oxidation at the same temperature. This was due to the faster oxidation rate
caused by oxide scale spalling associated with thermal cycling. The inner depletion zones were, however, similar for both thermal cycling and isothermal
oxidation at 1100 °C. For the higher temperatures, ≥ 1000 °C, β phase could be
observed to form in the substrate due to interdiffusion.
The model was largely successful in predicting the main features of microstructural development with high-temperature exposure and showed acceptable agreement with experimentally established through-coating compositional profiles. The β-depletion life of the coating was predicted to be ∼ 220 h
and ∼ 3000 h for 1100 °C and 1000 °C respectively. For 900 °C, the life was in
excess of 10000 h, which was the maximum simulation time used.
The work provided an important complement to fatigue-based life prediction approaches. Whereas a fracture mechanics-based approach may be useful
for thermal cycling, TGO growth and Al depletion based models may be more
appropriate for isothermal high-temperature exposure or thermal cycling with
very long hold times.
Paper VII: Microstructure-based life prediction of thermal barrier coatings
The research presented in this paper expands upon the β-depletion based model described in the paper Modeling of microstructural evolution and lifetime
prediction of MCrAlY coatings on nickel based superalloys during high temperature oxidation. The β-depletion based model was adapted to APS coatings
which also meant that internal oxidation had to be considered due to the formation of oxide stringers in the bond coat. The studied specimens were TBC
coated while the original model was developed for MCrAlY coated specimens.
The studied TBC system consisted of APS NiCoCrAlY and APS zirconia. The
specimens were isothermally oxidised up to 2500 h at 900 °C and 980 °C. The
specimens had not failed at the stop of the test and were still in very good condition. It could therefore only be concluded that the isothermal life of the coating was À 2500 h.
The TGO thickness was measured on cross-sections and used as input for
the β-depletion based model. It was found that the amount of internal oxi55
dation, which was measured by image analysis, remained fairly constant once
formed; the amount of internal oxidation was higher for higher temperatures.
Both interface TGO growth and internal TGO growth were assumed to follow a
power-law type equation (see Eq. 3) with n = 3 for interface TGO growth and
n = 5 for internal oxidation.
EDS was performed on a cross-sectioned as-sprayed specimen to get compositional profiles for input into simulations. The DICTRA software was used
to model the size of the inner and outer β-depletion zones. The model results agreed well with experimental observations for 980 °C but considerably
worse for 900 °C. The poor agreement between experimental and simulated βdepletion zone size occurred since the specimens oxidised at 900 °C developed
γ/γ0 zones adjacent to the γ/β zone; DICTRA failed to predict this and, as a
result, underestimated the size of the β-depletion zone.
A few life criteria were tested for the specimen oxidised at 980 °C; based on
literature, the end of the life was considered to have been reached when 1) complete β-depletion had occurred, 2) the TGO thickness reached 5 µm, or 3) the
Al content in the bond coat had dropped to 3 wt.%. β-depletion occurred at
∼ 600 h, the TGO-based criterion gave a life of ∼ 1450 h and the Al-depletion
based gave ∼ 2250 h based on simulation results and 3500–4200 h based on extrapolation from experimental data. The β-depletion based life criterion was
obviously too conservative, as was the TGO-based criterion. The Al-depletion
based life criterion appeared to give more reasonable results; however, the Al
depletion simulation in DICTRA overpredicted the Al depletion rate somewhat
and gave a more conservative life than the experimental Al-depletion indicated.
The study improved the previously developed β-depletion based model by
also including internal oxidation; internal oxidation can be very pronounced in
APS bond coats.
56
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