Volume 6 Preprint 72
Evaluation of the Critical Oxide Thickness to Initiate Spallation from a LPPS CoNiCrAlY Coating
S. Gray, M. P. Taylor, E. Chau and H. E. Evans
Keywords: LPPS, Overlay coatings, MCrAlY, TGO, critical oxide thickness, spallation
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Volume 6 Paper H016
Evaluation of the Critical Oxide Thickness to
Initiate Spallation from a LPPS CoNiCrAlY
S. Gray, M. P. Taylor, E. Chau and H. E. Evans
Department of Metallurgy and Materials, The University of Birmingham,
Edgbaston, Birmingham B15 2TT, UK, S.Gray@bham.ac.uk and
M.P.Taylor@bham.ac.uk and H.E.Evans@bham.ac.uk
The critical oxide thickness required to initiate spallation from an LPPS
CoNiCrAlY overlay coating during cooling from 1100°C to room
temperature has been determined metallographically and has been
found to lie within the range 2.5 to 3.2 µm. This finding compares
favourably with a finite element model developed for a compatible
alloy system in which delamination and spallation are assumed to be
due to the mechanism of wedge cracking at the TGO coating interface.
Keywords: LPPS, Overlay coatings, MCrAlY, TGO, critical oxide
Metallic overlay coatings of the MCrAlY type (where M is Ni, Co or a
combination of both) are regularly used in high-temperature plant to
provide resistance to oxidation and high-temperature corrosion
through the formation of a protective surface layer of alumina. Such
coatings are also frequently used as a bond coat in thermal barrier
coating systems. For both applications, the endurance of the coating
system depends on the maintenance of the protective alumina layer.
To enable this, a high resistance to spallation is required and also a
sufficient reservoir of aluminium within the coating so that re-healing
of the alumina layer is possible should spallation occur.
Experimental studies on a variety of alloy systems ranging from
chromia-forming austenitic steels [1,2], alumina-forming ferritic steels
[3,4] to alumina-forming nickel-based alloys  have shown that
spallation, initiated during cooling, requires a critical temperature drop
for a given oxide thickness. This critical condition can be understood
from the premise that oxide spallation will result when the strain
energy within the oxide layer equates to that required to produce
Assuming that the strain energy derives solely from
differential thermal contraction strains during cooling, it can be shown
ξE (α m − α ox ) (1 − υ )
Here, E and ν are, respectively, the Young’s modulus and Poisson’s
ratio of the oxide, αm and αox are, respectively, the coefficients of
thermal expansion of alloy and oxide, ξ is the oxide thickness, ∆Tc is
the critical temperature drop for spallation and γF is the effective
fracture energy for interfacial fracture.
Finite element analyses [6,7] have shown that this simple “strain
energy” criterion can be used to predict spallation arising from the
growth of a wedge crack along the oxide/metal interface.
important requirement, though, is that creep relaxation in the alloy, at
the crack tip, produces a period of zero crack growth during the early
stages of cooling and that final growth, in the vicinity of the
temperature drop, ∆Tc , occurs rapidly . Under these conditions, the
effective fracture energy, γF , can acquire high values, typically 10-100
J.m-2, depending on the level of strain energy dissipated by local
substrate alloy creep.
These concepts and the model of spallation will be compared, in this
paper, with the behaviour of a CoNiCrAlY overlay coating. Attention
will be paid to the morphology of the spall zones as a guide to the
nature of the spallation process.
The experimental results will be
compared with earlier  finite element calculations for a similar
NiCrAlY coating alloy.
supplied by Chromalloy UK and Alstom Power Ltd. in the form of
10mm diameter bars of 150 mm length, which had been recessed at
10mm intervals along the length and subsequently coated, Figure 1.
All corners had been chamfered prior to coating to avoid stress
concentration at the edges. The coating used was AMDRY 995 powder
deposited by Low Pressure Plasma Spraying (LPPS) to a thickness of
100 µm. Compositional details of the substrate and coating are given
in Tables 1 and 2, respectively.
Once coated, the specimens were
polished to produce an essentially flat outer surface. The bars were
partitioned for oxidation testing, producing cylinders of 10 mm
length, all exposed substrate surfaces were masked with an aluminabased paint prior to testing.
Schematic representation of the as-received specimen bar showing
dimensions and the mid-line along which sectioning of the specimens was
performed after testing.
Table 1 Nominal chemical composition of CM186-LC
Table 2 Nominal chemical composition of Amdry 995
The specimens were subjected to isothermal oxidation in laboratory air
at 1100°C for varying times up to 335 hours.
specimens were removed from the furnace at temperature and cooled
naturally in laboratory air to room temperature and left for a few
The cooling rate varied throughout this transient from an
estimated average value over the first 100oC of approximately 104 oC/h
to appreciably lower values at lower temperatures.
were visually inspected for signs of gross spallation.
In order to
maintain the integrity of the remaining oxide, each specimen was gold
sputtered and nickel plated in a Watts bath. The specimens were the
sectioned transversely at the mid-length using a slow-speed diamond
cutting wheel and mounted in cold setting resin. These were ground
on progressively finer SiC paper from 120 to 1200 grit and polished
progressively down to a ¼ µm diamond finish. Ultrasonic cleaning was
performed and finally sputtering with gold to eliminate charge build
up on analysis in the SEM.
The specimens were first examined using an optical microscope to
identify the approximate fraction of spallation that had occurred.
Subsequently a JEOL 6300 Scanning Electron Microscope, SEM, was
used to measure the oxide layer thickness distribution around the
specimen circumference and to determine the composition of the
oxide using Energy Dispersive Spectroscopy, EDS.
3.1 Specimen microstructure
The microstructural detail of the as-received specimen can be seen in
Figure 2. The SEM images reveal the duplex structure of the aluminium
Al-rich β phase (NiAl) and the γ′ phase (Ni3Al) present in the asreceived condition and also the presence of the α–chromium phase.
After time at temperature, a thermally grown oxide (TGO) layer
developed at the outer surface of the overlay coating, Figure 3, with a
reduction in the amount of β phase present in the coating in the
proximity of the surface. A few diffusionally isolated regions of the
overlay coating were identified in these specimens and discussed
elsewhere . In these regions, internal oxidation produces diffusion
barriers resulting in premature chemical failure of sections of coating.
For this present study, however, regions of the coating that have
experienced this diffusion cell behaviour have been avoided.
analysis confirms that in all cases the TGO included in this study was
SEM images of sections through the as-received specimens, (a) higher
magnification showing the phases present and limited internal oxidation and (b)
lower magnification showing the complete coating system.
Figure 3. SEM images of sections through a specimen held at 1100°C for 1hr showing
TGO formation on the outer surface and a reduction in amount of the darker NiAl
phase in the outer level of the coating associated with aluminium depletion at (a)
x800 and (b) x3500 magnification.
3.2 Oxide Growth Kinetics
TGO thickness measurements were made at randomly distributed sites
around the circumference of the oxidised specimens.
showing the frequency of distribution for 4 and 20hrs are given in
0.25 0.50 0.75 1.00 1.25 1.50 1.75 2.00 2.25 2.50 2.75 3.00
Figure 4. Histograms showing the distribution of TGO thickness in specimens held
at 1100°C for (a) 4hrs and (b) 20hrs.
It is clear from Figure 4 that the distribution of oxide thicknesses is
unimodal and, nominally, Gaussian.
The mean TGO thickness and
standard deviation taken from the specimens oxidised up to and
including 20hrs are presented in Figure 5(a). Significant amounts of
spallation of the TGO occurred on cooling to room temperature in
specimens oxidised for longer times but the thickness of the residual
TGO on these specimens was not included here. In Figure 5 (b) the
oxide thickness data are plotted logarithmically against time and
demonstrate a power law dependence (equation 2):
where x is the TGO thickness, t is time at temperature and kp is the
rate constant. From Figure 5(b) the value for n was found to be 2.5.
These kinetics are plotted on linear axes in Figure 6 to show that the
data extrapolate convincingly to zero time and, so, indicate that there
has been a negligibly short period of non-protective transient
oxidation. The rate constant Kp obtained from this figure is 1. 28x1019
log time (s)
TGO Thickness / µm
Time / hr
log TGO thickness
Figure 5. (a) Mean TGO thickness for specimens oxidised up to and including 20hr
at 1100°C and (b) the log time vs log thickness plot demonstrating an exponential
factor of 2.5 applies to this data.
TGO Thickness^2.5 / µm^2.5
Time / hr
Figure 6. Plot of the TGO thickness to the power of 2.5 against time. The slope of
this graph produces a value for the growth rate of 1.28x10-19 m2.5s-1.
Microscopical examination of cross sections of each of the specimens
was used to estimate the linear fraction of the specimen cross section
that had exhibited oxide spallation. These fractions, expressed as a
percentage are plotted in Figure 7 as a function of the exposure time
and can be used to identify the minimum exposure required to initiate
spall. The criteria for spallation here was the total loss of TGO from
the overlay coating. It can be seen that significant spallation occurred
in all specimens oxidised for times greater than 50h and that the
quantity increased with time at temperature and thus TGO thickness.
Significantly, for shorter times there was no detectable spallation, as
thickness was required to initiate spallation during cooling to room
From Figure 7, it can be seen that this critical oxide thickness is
developed during an exposure of between 20 and 50hrs at
A detailed SEM examination of the specimen oxidised
for 20hrs and 50hr was made. No evidence of spallation was found in
the 20hrs specimen, but at 50hrs, spalled regions of the order of 20µm
were found, Figure 8. It should also be noted from this figure that
there were no extensive regions of decohesion or void formation under
the oxide adjacent to these spall zones and that the oxide/metal
interface was sensibly planar
Percentage Cross Sectional Spallation
Time at Temperature/hr
Figure 7. Plot of the percentage of spallation as a function of time at temperature.
Ni plate TGO
Ni plate TGO
Figure 8. Two SEM images of sections through the LPPS CoNiCrAlY coating held at
1100°C for 50hr showing spallation sites. Note that the adherent oxide at the spall
site tends to have an inclined fracture surface and that there are no obvious defects
at the interface of this oxide with the underlying coating which remains sensibly
A more direct estimate of determining the spall initiation condition can
be obtained from Figure 9 which plots the percentage of spallation as
a function of the measured TGO thickness. This method uses all the
oxide thickness data and includes residual TGO thickness, i.e. where
spallation over some of the specimen surface will have occurred. It is
not possible to obtain a single, unambiguous value of the average
critical oxide thickness from this figure but a value in the range 2.5µm
to 3.2 µm seems reasonable. From the oxidation kinetics, equation (2),
this corresponds to exposure times of 22 and 40 hours respectively,
Percentage Cross Sectional Spallation
and is consistent with the estimate mad from figure 7.
Mean TGO Thickness/µm
Figure 9. Plot of the percentage of spallation as a function of mean
As can be appreciated from Figure 8, isolated regions of spallation
occur and these tend to be dispersed randomly across the specimen
surface with no obvious preference for edge locations. The examples
shown in Figure 8 represent the morphology after cooling from 50
hours exposure, i.e. shortly after spall initiation.
In this case, the
oxide thickness is approximately 3 µm (see preceding Section) and the
spall zone has an apparent diameter, 2r, of around 20 µm. This gives
a value of approximately 3 for the ratio r/ξ and this has profound
implications for the nature of the spallation process. In particular, it
can readily be shown that buckling of the oxide layer as a route to
spallation is not feasible. Thus, using Timoshenko’s  solution for a
biaxially stressed (in compression) clamped plate and assuming that
the only source of stress derives from differential thermal strains
during cooling, it can be shown  that the critical temperature drop,
∆Tb, to produce buckling is:
− α ox ) (1 − ν ) R
where the various terms have already been defined. Inserting typical
values into this equation of ν=0.27, ξ=3.0x10-6 m, R=10x-6 m and
(αm-αox)=9x10-6 K-1 gives a value for ∆Tb of approximately 13000K.
Clearly, this is far too large for buckling to be a feasible route to oxide
spallation but it should be emphasised that the calculation also
assumes that a large 20-µm zone of decohesion exists at the
oxide/metal interface prior to cooling.
In actuality, no significant
voidage, at a spatial resolution of around 1 µm, was detected at the
oxide/metal interface. Such a small zone of decohesion implies that
only very thin oxide could develop a buckle configuration during
cooling as shown by the left-hand curve of the spallation map of
Figure 10 . This line has been calculated using equation (3) with a
value of 1 µm for R and shows that buckling could occur in the present
tests only for oxides < 0.1 µm thick.
The critical oxide thickness for spallation deduced from the present
work lies in the range 2.5 to 3.2 µm but these are means and, as
shown in Figure 4, a distribution of values of oxide thickness exists.
Nevertheless, the low tail of this distribution at 20 hours exposure, i.e.
near the spall initiation condition, still provides oxide thickness values
far larger, at 2 µm, than those required to initiate buckling.
emphasised that this is the case for specimens measured after
exposures at which spallation had not occurred so that the full
distribution of oxide thicknesses was available for measurement.
can be concluded, therefore, that the spallation found in the present
tests is not caused by buckling of particularly thin regions of the oxide
A spallation map for the Ni16CrAlY system cooled from 1100°C at the
constant rates shown . The superimposed line is for the present CoNiCrAlY alloy
naturally cooled from 1100°C.
The cross sections of the spall zones shown in Figure 8 also show that
the fracture surface of the adjacent, adherent oxide tends to be
inclined at roughly 45o to the plane of the oxide/metal interface. This
is the geometry expected [1,6] when spallation results from the growth
of a wedge crack along the oxide/metal interface, as shown
schematically in Figure 11. Here, the inclined surfaces are the planes
over which shear cracking has occurred early in the cooling transient.
Subsequent sliding during further cooling then results in the
development of out-of-plane tensile stresses across the oxide/metal
interface and the growth of a wedge crack . It should be noted that
the sliding displacements required to develop the wedge crack are,
typically, very small , of order an elastic deflection, and can be
accommodated even though the inclined surface may not be
Figure 11. (a) Schematic representation of the formation of wedge cracks at the base
of shear cracks within the TGO due to thermal contraction differences between the
TGO and the coating. (b) Joining of wedge cracks leading to spallation.
As outlined in the Introduction, finite element predictions [7,8] show
that the wedging process can be described by equation (1) provided
that coating creep is accounted for. Energy dissipation by creep then
results in effective fracture energies larger than required to propagate
an elastic crack. These will increase as the extent of creep relaxation
increases, for example, as cooling rate reduces. This has been shown
 for a free-standing CoNiCrAlY coating, similar to the current,
during cooling from 11000oC where γF was found to be 22 J.m-2 at a
cooling rate of 104 oC/h but to increase to 40 J.m-2 for the slower
cooling rate of 102 oC/h.
The corresponding predictions of critical
temperature drop to initiate spallation, Equation (1), are shown as the
two right-hand curves in Figure 10.
It can be appreciated that unlike the buckling route, oxide spallation
by wedging becomes energetically easier with thicker oxide layers. For
the case shown in figure 10, however, there is a separation of the
buckling and wedging lines so that, at these intermediate oxide
thicknesses, the oxide layer remains adherent and mechanically stable
during cooling from 1100oC. The fact that the buckling and wedging
lines do not intersect on this map means that any buckle formed, e.g.
with thin oxides, will not propagate laterally .
The faster cooling rate of the two used in constructing the wedging
lines shown in Figure 10 is a reasonable approximation to the
experimental cooling rate experienced in the present tests during the
early stages of cooling.
The lower of the two (102 oC/h) is more
representative of that occurring at intermediate temperatures during
the transient. Even so, although not modelling accurately the variation
in actual cooling rates through the transient, these previously
computed lines will still offer guidance on the value of the critical
oxide thickness required to initiate spallation.
The range of values
(2.5 to 3.2 µm) deduced experimentally are shown as two points on
Figure 10 corresponding to a temperature drop of 1080oC. The
position of these within the region where wedge cracking is predicted
and close to the border of the stable oxide region is in broad
agreement with the model and consistent with the spall morphology
found in this study.
A study has been made of the oxidation and spallation behaviour of a
LPPS CoNiCrAlY overlay coating on a Ni-based CM186-LC superalloy
after oxidation in air at 1100oC. It is found that the growth of the
protective alumina layer is sub-parabolic, characterised by a time
exponent of 0.4.
Oxide spallation occurred during cooling to room
temperature after an exposure time of between 20 and 40 hours,
corresponding to a critical oxide thickness in the range 2.5 to 3.2 µm.
SEM examination showed that the spall sites were, typically, around 20
µm in diameter and tended to show inclined fracture surfaces. It was
demonstrated that a buckling route to spallation could not operate and
that an oxide wedging process led to the observed spallation.
observed values of the critical oxide thickness to initiate spallation
were in broad agreement with earlier finite-element analyses on a
free-standing CoNiCrAlY coating alloy.
The authors would like to thank Mr. Mick Whitehurst and Mr. Mike
Henderson of ALSTOM Power (Ltd) and Mr. Rodney Wing of Chromalloy
United Kingdom Ltd. for the provision of specimens and their
encouragement during the course of this work. The Engineering and
Physical Sciences Research Council fund the project under grant
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