Volume 6 Preprint 68
Delamination Processes in Thermal Barrier Coating Systems
H.E. Evans and M.P.Taylor
Keywords: TBC lifing, oxidation, chemical failure
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Volume 6 Paper H011
Delamination Processes in Thermal Barrier
H.E. Evans and M.P.Taylor
Metallurgy and Materials, University of Birmingham, Edgbaston,
Birmingham, B15 2TT, UK, H.E.Evans@bham.ac.uk and
Thermal barrier coatings are used widely in both aeroengines and
land-based gas turbines but realisation of their full potential remains
hampered by incomplete understanding of the processes leading to
delamination and the spallation of the ceramic top coat.
development of cracks either at the interface of the thermally grown
oxide (TGO) with the bond coat or within the top coat close to or at its
interface with the TGO. Mechanisms that could result in such cracks
are considered in this paper.
Particular attention is paid to the
consequences of the growth of the TGO on the bond coat surface since
this takes place in a constrained environment and the associated
volume changes must be accommodated within the TBC system. An
important aspect is the topology of the bond coat surface. It is shown
how out-of-plane continuity strains favouring delamination can
develop isothermally at the oxidation temperature on non-planar
surfaces even when the oxide layer remains α-alumina and protective.
The situation will be exacerbated if aluminium depletion is sufficient
locally to trigger chemical failure and the formation of faster-growing
Further propagation of continuity cracks will be
favoured during cooling as a consequence of out-of-plane tensile
stresses resulting from differential thermal contraction strains.
Finally, finite-element results are provided of the growth kinetics of a
wedge crack along the TGO/bond coat interface and the sources of
stored energy that drive the crack are discussed.
Keywords: TBC lifing, oxidation, chemical failure.
Thermal Barrier Coatings (TBC) systems have the potential of offering
an improvement in gas turbine efficiency by increasing the inlet
temperature and reducing the amount of cooling air required in hightemperature components [1,2]. TBC systems are designed to confer
oxidation resistance through the formation in service of a protective
thermally grown oxide (TGO), usually alumina. This forms on a bond
coat, typically PtAl or MCrAlY (where M is Ni or Co or a combination of
both), an intermediate layer between the superalloy substrate and an
outer thermal barrier. The outer layer is usually a ceramic, generally
ZrO2/8%Y2O3 having low thermal conductivity. In a heat-flux situation,
e.g. with internal cooling, a component can then operate at
temperatures substantially less, e.g. 100-200oC, than that of the outer
surface of the ceramic coating. In highly-rated situations, these outer
temperatures would lead to rapid oxidation of the bond coat and
failure of the coating system. The mechanical integrity of the outer
ceramic coating is then, clearly, an important issue.
understanding of the mechanisms of the delamination processes, has
meant, however, that TBCs tend to be used in a conservative manner.
This approach ensures that the component will have an adequate life
even in the absence of the coating system; the purpose of the TBC is
then simply to extend service lifetimes [1-3]. TBCs have been used for
some years on turbine combustors and after burners.
these areas to apply TBC systems routinely to nozzle guide vanes and
aerofoils is necessary to realise their full economic potential. This is
hampered by limited understanding of the failure mechanisms and the
consequent inability to predict coating lifetime reliably (lifing).
This paper considers possible delamination processes in TBC systems.
It is recognised that the final catastrophic spallation of the ceramic top
coat will occur during the last fatal cooling transient, possibly by
Classic elastic theory  can be used to show that an
extensive zone of decohesion must exist at the start of cooling for
buckling to occur.
This zone will develop progressively during
exposure and the rate of its development will determine the life of the
coating. Indeed, it may be possible to use luminescence techniques
 to estimate TBC residual lifetime through monitoring the rate of
development of fracture damage. Accordingly, particular attention will
be paid in this paper to possible mechanisms that could lead to the
nucleation and slow growth of sub-critical delamination cracks. It will
be argued that the surface topology of the bond coat can be a critical
factor and that cracks can arise isothermally as a result of the growth
of the TGO in a mechanically-constrained situation. Crack growth by
wedging along flat TGO/bond coat interfaces will also be considered
as will the consequences of locally enhanced aluminium depletion.
The tendency over the last decade or so has been for TBCs with
MCrAlY bond coats to be used in industrial turbines. These coats are
usually deposited by plasma spraying, initially by air plasma spraying
(APS) or with an argon gas shroud (ASPS) but later using high velocity
electroplating. The early APS or ASPS coatings were of relatively low
density and permitted ingress of the oxidant along the boundaries
between the splat particles.
The consequent internal oxidation
resulted in the diffusional isolation of some splat particles and the
formation of diffusion cells .
The rapid depletion of aluminium
within these led to rapid failure, termed chemical failure, through the
formation of non-protective, Ni-rich oxides.
High density coatings,
e.g. those manufactured by LPPS, have been found  to exhibit
substantially reduced diffusion cell formation. In all these cases, the
outer ceramic coating is usually deposited by APS for which a relatively
rough bond coat surface is required to aid mechanical keying between
the two layers. An example of a TBC system consisting of a NiCrAlY
bond coat, an APS top coat and a Nimonic 80A alloy substrate is
shown in Figure 1 . The horizontally layered structure of the splat
boundaries formed by plasma spraying is particularly obvious in the
Figure 1: SEM of an as-sprayed TBC system.
TBCs with PtAl bond coats, produced by diffusional aluminising, tend
to be used in aeroengine applications with a top coat deposited by
electron beam vapour deposition (EBPVD).
An example is given in
Figure 2. In this case, the deposition process results in a columnar
structure to the top coat in which the bonding between the vertical
columns is relatively poor. The bond coat tends to be oxidised prior
to top coat deposition to aid adherence and to obviate the need for as
rough a bond coat surface. In practice, though, there can be a large
batch-to-batch variation in the bond coat surface topology.
Another feature of this processing route, as can be seen in Figure 2, is
the presence of an extensive interdiffusion zone between bond coat
and alloy substrate. This indicates that there is no significant barrier
to solid state diffusion across the interface with the alloy substrate and
temperature exposure. In principle, this chemical contamination could
affect the adherence of the thermally grown oxide (TGO) layer or of the
top coat [9, 10]. By contrast, plasma-sprayed MCrAlY coatings show
much less interdiffusion (Figure 1) indicating that some barrier to
diffusion does form between the bond coat and the alloy substrate.
Even so, diffusion of elements, particularly titanium, can still occur
during high temperature exposure and have been identified  at
oxide spallation sites of overlay coatings in the absence of a top coat.
Figure 2: Back scattered image of a section through a platinum
aluminide bond coat with a EBPVD YSZ topcoat which has been held at
1100°C for 1h.
A Simple Lifing Model Based on TGO Thickness
Figure 3 provides a compilation of literature data [8, 10, 12-21] of
times to failure as a function of temperature of both main types of TBC
systems. This plot includes results from tests cycled with frequencies
ranging from isothermal exposure to those using less than hourly
cycles for which the accumulated time at peak temperature is used. A
distinction has been made on the diagram between these various types
of specimens and test procedures.
Time to Spallation
MCrAlY bond coat with APS topcoat
MCrAlY bond coat with APS topcoat
MCrAlY bond coat with APS topcoat
MCrAlY bond coat with EBPVD topcoat
MCrAlY bond coat with APS topcoat
MCrAlY bond coat with APS topcoat
MCrAlY bond coat with EBPVD topcoat
MCrAlY bond coat with APS topcoat
Pt-aluminide bond coat with EBPVD topcoat
Pt-aluminide bond coat with EBPVD topcoat
Pt-aluminide bond coat with EBPVD topcoat
Pt-aluminide bond coat with EBPVD topcoat
ξ = 8µm
ξ = 5µm
ξ = 3µm
Figure 3: A compilation of literature data [8, 10, 12-21] showing the
dependence of TBC lifetimes on oxidation temperature. Both of the
main types of TBC systems are included together with test conditions
ranging from isothermal to frequent thermal cycling.
represent the time at temperature required to produce a TGO of the
This figure shows that not only is there an obvious reduction in life
with increasing temperature but that there does not appear to be a
systematic trend with specimen type or test procedure.
This is a
surprising result but one that must be viewed with caution because of
the large scatter between different sets of data and the limited number
of results in any one set. It should also be noted that the data shown
are for current production quality specimens, as described in the
preceding section, and exclude those modified to increase life. It is
known, for example, that Pt additions can effect such an improvement
[9,18,22] as can aluminising  and pre-oxidation of bond coats
The trend of decreasing TBC lifetime with increasing temperature,
shown in Figure 3, is reasonably similar to that expected if failure
occurred at a critical thickness of the TGO layer, independent of test
This is indicated in the figure by the lines which
correspond to a mid-range critical TGO (alumina) thickness, ξ, of 5 µm
with approximate lower and upper bounding values of 3 µm and 8 µm,
These values were calculated assuming pseudo-
parabolic growth of the alumina layer with a rate constant, kp, given, in
= 2.0 × 10 − 6 exp
where t is the oxidation time in seconds and T the temperature in K.
This equation derives from the compilation of Hindum and Whittle 
and corresponds to their more oxidation-resistant alumina-forming
alloys over the temperature range 1000-1200oC. These well-behaving
alloys tend, in fact, to be bond coat materials, either MCrAlY  or
PtAl . The predictions of equation (1) agree well enough with the
few measurements of oxide thickness obtainable from the data
sources used to produce Figure 3.
This simple lifing model of failure at a critical TGO thickness correctly
emphasises the importance of the formation of this on TBC endurance.
It is a concept that has been recognised, in principle, for many years
but, nevertheless, understanding of the detailed mechanisms by which
oxide growth promotes delamination remains elusive.
As can be
development prior to gross spallation of the top coat, cracking in the
top coat (marked A) or where local non-protective oxides (chemical
failure) have formed (marked B) tends to be above the TGO within the
top coat. The TGO can then, in no sense, be said to have “weakened”
the bond coat/top coat interface nor will its stored energy, resulting
from growth and thermal stresses, contribute to the top coat cracking.
However, it is also clear from Figure 4 that cracking can occur at the
TGO/bond coat interface (marked C). In this case, a contribution to
the driving force for cracking will derive from the release of oxide
The importance of bond coat surface roughness,
produced as-manufactured [11, 26-29], resulting from local chemical
failure [6,27] or caused by thermal cycling  is now recognised. It is
likely to be a significant, and largely uncontrolled, factor contributing
to the factor 10 spread in TBC lifetimes evident in Figure 3. Various
aspects of the influence of bond coat surface roughness will be
addressed in the remainder of this paper.
Figure 4: Micrograph of a section through a TBC with an electroplated
CoNiCrAlY bond coat and an APS YSZ topcoat held isothermally at
1100°C for 600h showing cracking within the top coat (marked A),
local chemical failure (marked B) and cracking at the TGO/bond coat
interface (marked C).
Constrained Oxide Growth on a Rough Bond Coat Surface
The key to understanding the important influence of the topology of
the bond coat interface is the recognition that the TGO grows in a
mechanically constrained environment. This is determined largely by
the elastic properties of the top coat but modified by the creep
properties of the bond coat. Stresses will develop in an out-of-plane
sense within the top coat whenever there is a variation of upward
displacement rates along the bond coat surface. These may arise from
differences in oxidation growth rates, as discussed previously , but
can also arise simply from geometrical effects even when the intrinsic
oxidation rates are everywhere constant.
It is also important to
recognise that this process may lead to cracking (termed continuity
cracking in this paper) at the oxidation temperature under isothermal
The importance of the constraint offered by the top coat has been
recognised for some time, albeit implicitly, in finite element codes that
model the growth of the oxide layer on a rough bond coat surface at
temperature [31-34]. In all these cases, though, temperature cycles
were also imposed and these, together with the considerable
complexity within the code, makes it difficulty to appreciate the
physical significance of the processes that would have occurred during
isothermal exposure at the oxidation temperature.
To aid in this
visualisation, consider oxidation of a bond coat having a sinusoidal
profile as shown schematically in Figure 5.
The direction of TGO
growth is perpendicular to the localised bond coat surface so that at
peaks and troughs or on a planar surface the direction of growth is
However, at any point away from these the growth
direction will have a vertical and horizontal component and the former
will have a magnitude less than that at the peaks and troughs. This
will produce a variation in the upward displacement rates along the
bond coat surface, Figure 5.
The imposed displacements produce
continuity strains and associated tensile stresses within the top coat,
as shown schematically in Figure 6. The term “tensile wings” has been
coined to describe the pattern of stresses developing at temperature in
Figure 5: Oxidation of a rough bond coat produces a variation in
upward displacement rates along the surface.
Modelling estimates of the magnitude of these stresses will not be
reliable unless creep in the bond coat is allowed for. This is a crucially
important factor in the modelling process but because creep rates will
be non-linear in both stress and temperature, numerical (finite
element) procedures rather than analytical need to be used.
tensile stresses prove to be sufficiently high, then cracks would be
expected to form on the flanks of the hills of the bond coat surface
and some evidence for this is given in Figure 7.
Even though no
reliable estimates of the magnitude of these stresses have yet been
produced, it is intuitively reasonable to expect them to increase with
the roughness ratio α⁄λ, where α is the amplitude of the sine wave
representing the surface topology of the bond coat and λ is the
Figure 6: The imposed displacements produce continuity strains and
associated tensile stresses within the top coat. In the simplest case of
a very rough interface these could manifest as “tensile wings” along
the flanks of the bond coat protuberances, as shown schematically.
At the oxidation temperature, the top coat above the regions of
maximum upward displacement, in particular above the peaks in the
bond coat, experience out-of-plane compressive stresses. However,
during cooling these can become tensile due to the differential
thermal contraction strains and can then aid the lateral growth of the
continuity cracks produced at the oxidation temperature. It is by this
means that a large delamination crack can develop within the top coat
during repeated thermal cycles and periods of high temperature
Figure 7: Back scatter SEM image of section through a TBC consisting
of a Pt aluminide bond coat with an EBPVD YSZ top coat that has been
held at 1200°C for 3h, showing cracking (arrowed) in the top coat
associated with the tensile wings.
Chemical failure arises when insufficient aluminium remains in the
bond coat to maintain a protective alumina layer. A framework theory
for the initiation of chemical failure has been developed elsewhere 
and applied both to overlay  and to TBC systems with APS bond
As a result of chemical failure, the TGOs formed at this
stage in the life of the coating are non-protective, grow quickly and
cause rapid deterioration of the coating. Another major factor is that
chemical failure is likely to occur initially in discrete localised regions
and, so, will result in an extreme variation in oxide growth rates across
the bond coat surface. This process has been studied in some depth
by Evans and Taylor  for the case of an APS bond coat.
discrete volumes of the coating were found to become diffusionally
isolated (diffusion cells) and to suffer rapid aluminium depletion and
early chemical failure. The process is shown schematically in Figure 8
from which it can be appreciated that the spatial variation in upward
displacement rates along the bond coat surface will lead to out-ofplane tensile stress development within the top coat but located over
regions which are still oxidising at low rates.
In practice, relatively
low-density APS bond coats are now seldom used and this type of
failure, associated with the extensive formation of diffusion cells
within the bond coat, is unlikely to be an important issue with modern,
dense bond coats.
Out of plane tensile stresses
Break away oxidation
Figure 8: Spatial variations in the thickness of oxides across the
surface of the bond coat leads to out of plane stresses developing
between such areas and the possible formation of delamination cracks
in these regions.
More generally though, and of application to dense MCrAlY or PtAl
bond coats, any spatial variation of TGO growth rates, whether
protective or breakaway, will lead to out-of-plane stresses at the
oxidation temperature. Their magnitude will, amongst other things,
depend on the spread of oxide growth rates and, obviously, it will be
of benefit to have a narrow distribution of these.
Localised chemical failure can also arise in regions where, due to the
local geometry, aluminium depletion by alumina formation is greater
that the rate of aluminium replenishment.
This tends to occur in
regions with large ratios of surface area to volume, for example at
peaks in the bond coat surface, as shown schematically in Figure 9.
Analytical calculations of this balance of fluxes for the realistic case of
concurrent oxidation cannot be made.
However, by envisaging the
protuberance to be approximated by a sphere, it is expected, for this
purpose of illustration, that the time to aluminium depletion and
chemical failure at the protuberance tip will scale roughly inversely
with the roughness parameter, i.e. ∝ λ/α.
local Chemical Failure leads to
faster-growing Ni/Cr-rich oxides
Figure 9: Enhanced local Al depletion arises when Jox > JAl and may
lead to Chemical Failure and the formation of Cr/Ni-rich oxides at
bond coat protuberances.
A consequence of chemical failure in these regions is that fastgrowing, non-alumina oxides will develop at the protuberance tips, as
shown schematically in Figure 9.
As a result, a spatial variation of
upward displacement rates will extend across the bond coat surface in
an analogous manner to that described above for chemical failure of
flat APS bond coats. Unlike that case, however, this damaging process
does not arise from the deposition of a low-quality porous coating but
from the roughness applied to the surface of even a dense coating
during manufacture of the TBC system. Out-of-plane tensile stresses
will again develop isothermally at the oxidation temperature and it is
expected that these will peak above the troughs in the roughness
profile, i.e. between the bond coat protuberances. Estimates of the
magnitude of these stresses cannot be made reliably without the use
of finite element procedures that fully account for relaxation effects
due to bond coat creep but it can be appreciated from Figure 10 
that they can be sufficient to nucleate delamination cracking.
Extension of such cracks through the non-protective oxides overlying
the bond coats peaks can again occur during cooling when these
become the sites of out-of-plane tensile stresses. The process will be
helped by the presence of porosity in these non-protective oxides
Figure 10: Crack nucleation within the top coat between bond coat
protuberances undergoing chemical failure .
6. Failure of TBCs with Flat Bond Coats
The micrographs shown in this paper and the results contained in
Figure 3 have been from samples with bond coat topographies
representative of current commercial TBC systems.
In all cases,
relatively rough surfaces have existed and the tendency has been for
delamination cracking to develop in the top coat by one or other of the
mechanisms described above. For APS top coats, it has been found
necessary for the bond coat surface to be rough to aid mechanical
keying but this is not a pre-requisite with EBPVD top coats. For flat
surfaces and in the absence of edge effects, it is not clear how
delamination cracking can develop. One possibility, of course, is that
localised chemical failure will lead to the development of out-of-plane
stresses at the exposure temperature, as described above. It is known
, though, that good quality dense bond coats with a flat surface do
not readily enter chemical failure.
A related process, as indicated
above, is that non-uniform growth rates of the protective alumina
layer over the bond coat surface will also lead to out-of-plane stresses
at the oxidation temperature. This, undoubtedly, will be a factor but it
is, as yet, unclear whether significant stresses will develop.
Similarly, during cooling, flat interfaces will not experience out-ofplane tensile stresses unless there is some outward but localised
displacement of the surface layer . As has already been indicated,
buckling will not occur until an extensive delamination zone has
formed. Wedge cracking can occur, however, on planar interfaces and
can progressively develop zones of decohesion. An example of wedge
cracking is shown at C in Figure 4 to demonstrate that the process can
also occur on interfaces which, whilst not particularly convoluted, are
not geometrically flat.
Significantly, the crack path in this case lies
along the TGO/bond coat interface.
It has been noted previously
[9,29] that the crack path tends to be along the TGO/bond coat
interface for flat surfaces, in contrast to more convoluted surfaces
where, as discussed above, the fracture path tends to lie within the top
The mechanism of wedge cracking has been discussed
elsewhere in detail and applied to the conventional spallation of
protective oxide layers [39-43]. It envisages that the compressive inplane stresses developed in the TGO during cooling result initially in
shear failure of the oxide layer. Subsequent cooling permits sliding
across the interface, i.e. a wedging process, the development of outof-plane tensile stresses across the TGO/bond coat interface and the
growth of a delamination crack. These crack features can be seen at C
in Figure 4.
due to shear
Tensile stresses ahead
of tip of17
t f TGO
axisymmetric finite element model showing the locations of tensile
and compressive stresses.
Extensive finite element (FE) modelling of the wedge cracking process
in conventional alloys has now been performed [e.g. 40-42] but has
not previously been undertaken for a TBC system. The geometry of
the FE model used here is shown in Figure 11. It shows the central
part of an axisymmetric, multi-layered TBC system consisting of a
Haynes 230 base alloy, a 100µm thick MCrAlY bondcoat, a 0.1-µm
thick TGO/bond coat interfacial zone with the properties of alumina, a
2.5-µm thick alumina TGO layer and, finally, a 250µm thick partiallystabilised zirconia top coat. Further details of the FE approach have
been given elsewhere .
Interfacial finite elements are used to
model a pre-existing, 45o-inclined shear crack within the TGO and the
TGO/bond coat interfacial zone. The wedge crack is confined to this
latter zone and propagates when the out-of-plane tensile stress ahead
of the crack exceeds a critical value of 1700 MPa. The interfacial zone
is considered to behave elastically and, so, this critical stress
corresponds to a critical elastic fracture strain of approximately 0.4%,
consistent with literature values . This type of fracture criterion
rather than an energy balance is preferred in this context since much
of the stored energy is dissipated into bond coat creep rather than in
extending the crack.
The creep rate of the bond coat and the
substrate alloy were expressed as:
ε = Aσ n exp −
where A=3.237x107, n=3.0, Q= -298000 for the MCrAlY coating 
and A=5.915x102, n=6.885, Q= -410013 for the Haynes 230 alloy. σ
is the stress in Pa and T is absolute temperature.
parameters are given in Table 1.
Table 1: Materials Parameters used in the FE Modelling
x 106 K-
Zirconia Top Coat
1 - 90
The predicted kinetics of wedge crack growth during cooling from
1100oC at a constant rate of 40x103 oC h-1 are shown in Figure 12 for
three values of the Young’s modulus of the top coat. This approach
was used to establish the importance of the physical constraint of the
top coat on the wedging process and of its elastic stored energy on
propagation of the wedge crack.
Length of interfacial crack/µm
Temperature drop from 1100°C / °C
Young's modulus of top coat 90GPa
Young's modulus of top coat 10GPa
Young's modulus of top coat 1GPa
Figure 12: The kinetics of wedge crack growth during cooling the TBC
system from 1100oC at a constant rate of 40x103 oC h-1. The different
curves correspond to different assumed values of the Young’s
modulus of the top coat.
The first important point to note from these results is that wedge
cracking will be expected, at this TGO thickness of 2.5 µm, during
axisymmetric model used are relatively short (8 µm) radius, however
and it is unclear whether much longer cracks could develop from a
single wedging location. Nevertheless, some delamination cracking is
certainly expected, even though top coat spallation may not occur, and
this is consistent with the lower bound behaviour shown in Figure 3. A
second point to note is that the constraint of the top coat, as reflected
in its Young’s modulus, does affect the ease of nucleation of the
wedge crack. It can be seen that a temperature drop of around 250oC
is required with a (realistic) top coat modulus of 90 GPa but that this
drops to under 100oC when the modulus is reduced to 10 GPa. This
constraint arises even though the shear displacement on the shear
crack within the oxide layer is generally < 0.1 µm.
After nucleation, further growth of the wedge crack is inhibited by
stress relaxation at its tip by bond coat creep during cooling. This
result has been previously found with similar computations on
conventional alloys [40-42] and is another illustration of the
importance of incorporating creep relaxation processes in such
Propagation of the wedge crack continues (Figure 12) at
temperatures where creep relaxation becomes slow but it is of
particular interest to note that the rate of growth of the crack at these
lower temperatures is not sensitive to the Young’s modulus of the top
coat. This means that the elastic stored energy within the top coat
does not make a significant contribution to the driving force for wedge
crack growth along the TGO/bond coat interface in this system. This
is obtained from release of the stored energy within the TGO.
The lifetime of TBC systems of commercial quality reduces drastically
with increasing temperature with a dependence which is broadly that
for the growth of an alumina TGO between the bond coat and the
partially-stabilised zirconia top coat. At any given temperature, there
is a large scatter in the results, typically a factor 10 in lifetime, but
there does not appear to be a significant difference between
isothermal and thermally-cycled specimens. The inference is that time
at temperature, and the growth of a bond coat TGO, is the dominant
parameter that results in TBC degradation. The alumina thicknesses at
failure that reasonably bound the experimental data are 3 µm and 8
µm with a median value of 5 µm.
An important factor contributing to the large spread in TBC lifetimes is
thought to be variations in the roughness of the bond coat surface and
various possible mechanisms that lead to the nucleation and growth of
delamination cracks have been considered in this paper.
summarised in Figure 13 where the roughness parameter used is the
ratio of amplitude, α, to wavelength, λ.
wedge cracking along BC interface
Time to Failure
cracking in topcoat due
to continuity strains
continuity cracking in TC
enhanced by chemical
failure of BC
Increasing Roughness, α/λ
Figure 13: Schematic TBC Delamination Map relating lifetimes to bond
coat surface roughness.
The presence of bond coat surface roughness leads to a spatial
variation in continuity strains which will result in out-of-plane tensile
stresses at the exposure temperature. This will be made worse by the
development of fast-growing non-protective oxides due to aluminium
depletion (chemical failure) at the tips of bond coat protuberances as
roughness increases. Modelling of these delamination processes must
incorporate creep relaxation processes within the bond coat in order
to produce reliable estimates of stress. Progress in this area is being
sophisticated codes that permit concurrent oxide growth.
rough interfaces, it is expected that delamination cracks will develop
within the ceramic top coat.
With flatter interfaces, delamination can proceed by the growth of
wedge cracks which develop along the TGO/bond coat interface during
cooling. Finite element computations of this process are presented in
this paper. It is shown that the mechanical constraint of the top coat
can inhibit, but not prevent, crack nucleation for realistic values of the
Young’s modulus of the top coat. Creep relaxation within the bond
coat will then inhibit further growth until low temperatures are
achieved. It is found that the driving force for growth in this stage is
the release of energy stored within the TGO and that that within the
top coat does not contribute significantly.
We are grateful to the Engineering and Physical Sciences Research
Council for contributing funding to this work.
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