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The University of Manchester Research
A mechanistic understanding on rumpling of a NiCoCrAlYbond coat for thermal barrier coating applicationsDOI:10.1016/j.actamat.2017.02.003
Document VersionAccepted author manuscript
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Citation for published version (APA):Chen, Y., Zhao, X., Bai, M., Yang, L., Li, C., Wang, Y. L., Carr, J., & Xiao, P. (2017). A mechanistic understandingon rumpling of a NiCoCrAlY bond coat for thermal barrier coating applications. Acta Materialia, 128, 31-42.https://doi.org/10.1016/j.actamat.2017.02.003
Published in:Acta Materialia
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https://doi.org/10.1016/j.actamat.2017.02.003https://www.research.manchester.ac.uk/portal/en/publications/a-mechanistic-understanding-on-rumpling-of-a-nicocraly-bond-coat-for-thermal-barrier-coating-applications(95359148-6362-42bf-889c-2822d605c27b).html/portal/ying.chen-2.html/portal/p.xiao.htmlhttps://www.research.manchester.ac.uk/portal/en/publications/a-mechanistic-understanding-on-rumpling-of-a-nicocraly-bond-coat-for-thermal-barrier-coating-applications(95359148-6362-42bf-889c-2822d605c27b).htmlhttps://www.research.manchester.ac.uk/portal/en/publications/a-mechanistic-understanding-on-rumpling-of-a-nicocraly-bond-coat-for-thermal-barrier-coating-applications(95359148-6362-42bf-889c-2822d605c27b).htmlhttps://doi.org/10.1016/j.actamat.2017.02.003
1
A mechanistic understanding on rumpling of a NiCoCrAlY bond coat for thermal barrier coating applications
Y. Chen1, X. Zhao2, M. Bai1, L. Yang2, C. Li1, L. Wang1, J. A. Carr1, P. Xiao 1, 2†
1 School of Materials, University of Manchester, Manchester M13 9PL, UK
2 Shanghai Key Laboratory of Advanced High-Temperature Materials and Precision Forming,
Shanghai Jiao Tong University, Shanghai, 200240, PR China
Abstract
We present a series of experimental observations on surface rumpling of an initially flat NiCoCrAlY
coating deposited on a Ni-based superalloy during cyclic oxidation at 1150 °C. The extent of rumpling
of the coating depends on the thermal history, coating thickness and exposed atmosphere. While the
coating surface progressively roughens with cyclic oxidation, the bulk NiCoCrAlY alloys with the same
nominal composition are much less susceptible to rumpling under the same oxidation conditions. The
coatings, especially the thin ones, experience substantial degradation (e.g. β to γ phase transformation)
induced by oxidation and coating/substrate interdiffusion. The observations together suggest that
rumpling of the NiCoCrAlY coating is driven by a combination of the lateral growth of the thermally
grown oxide and coating/substrate thermal mismatch. The results in this work are further discussed
and compared with the rumpling behaviour of a β-(Ni,Pt)Al bond coat reported in the literature to
illustrate the importance of possible factors in governing the development of rumpling in the
NiCoCrAlY coating.
Keywords: Rumpling; NiCoCrAlY; Oxidation; Microstructure
1. Introduction
Thermal barrier coatings (TBCs) are used in hot sections of gas-turbine engines to protect the
superalloy components and improve the engine efficiency [1-3]. A typical TBC system consists of a
ceramic topcoat (typically made of 7-8 wt. % yttria-stabilised zirconia) and an intermediate metallic
bond coat (made of β-(Ni,Pt)Al, NiCoCrAlY or Pt-diffused γ/γ’ alloys) deposited on a superalloy
substrate. Upon exposure at operational temperatures, a thermally grown oxide (TGO), predominately
α-Al2O3, develops between the bond coat and the topcoat. The topcoat gives thermal insulation; the
† Corresponding author: P. Xiao ([email protected]).
mailto:[email protected]
2
TGO provides oxidation resistance; the bond coat promotes the formation of the protective TGO and
enhances bonding to the substrate; the superalloy bears the loads. The whole TBC system is dynamic
and all the components interact with each other to control the performance and durability of the TBCs.
One of the most important forms of TBC degradation is associated with progressive roughening, also
termed rumpling, of the bond coat surface along with the TGO scale during the course of cyclic
oxidation [4-9]. The out-of-plane displacement accompanying rumpling gives rise to a tensile stress
around the undulation peaks and a compressive stress around the undulation valleys, respectively, across
the TGO/bond coat interface [5], where the TGO scale is compressively stressed in the lateral
direction. The tensile stress can initiate interfacial cracking and eventually leads to failure of the TBC.
Surface rumpling of β-(Ni,Pt)Al coatings induced by cyclic oxidation has been extensively investigated
by experiments and simulations in past decades. These studies have shown that the evolution of
rumpling depends on oxidation time, oxidation temperature, oxidation mode (cyclic or isothermal),
oxidation atmosphere (in air or vacuum), sample configuration (bulk alloys or coatings), coating
thickness, content of reactive elements (e.g. yttrium and hafnium) in the coatings and external load [10-
16]. Balint and Hutchinson have developed a comprehensive analytical model (Balint-Hutchinson
model) to simulate rumpling development of the β-(Ni,Pt)Al bond coat during thermal exposure [17].
The mechanism of rumpling growth in the model is bond coat creep driven by the laterally compressed
TGO, facilitated by both thermal and stress activation of the creep mechanism, the lateral resulting
from strain mismatch between the bond coat and substrate and the phase transformation of the bond
coat. The predictions of the model have been found to be in agreement with most of the experimental
observations reported in the literature, although the correspondence between the experiments and
simulations remains to be validated.
Another type of bond coat, made of NiCoCrAlY alloys, also shows a propensity to rumple after
exposure at high temperature [18-23], although the coating shows other concomitant degradation
mechanisms apart from rumpling (e.g. thickness heterogeneities, often referred to as “pegs”,
accompanied by localised TGO cracks [23-27]). However, few experimental studies have been carried
out to systematically investigate rumpling of NiCoCrAlY coatings and how their rumpling could be
affected by experimental variables. The objective of this work, therefore, is to provide a systematic
3
study of rumpling of NiCoCrAlY coatings. The materials considered in this study include a
NiCoCrAlY bond coat and a NiCoCrAlY bulk alloy with the same nominal chemical composition.
Surface rumpling of these materials under a variety of experimental conditions is examined, together
with the associated TGO stress development and microstructural evolution. The observations are
further discussed and compared with the rumpling behaviour of a β-(Ni,Pt)Al bond coat reported in
the literature to illustrate the rumpling mechanism of the NiCoCrAlY coating and the importance of
possible factors in controlling the rumpling development of the NiCoCrAlY coating.
2. Materials and methods
2.1 Sample preparation
The NiCoCrAlY overlay bond coat was sprayed on one side of Hastelloy® X polycrystalline superalloy
strips (55× 8 × 1.6mm3) using high velocity oxygen fuel (HVOF) spraying. Table 1 shows the nominal
compositions of the as-sprayed coating and the Hastelloy® X superalloy substrate. The thickness of
the as-deposited coating varies between 200 and 230 μm from place to place (Fig.1a), owing to the
extensive surface roughness formed during the spraying process. The coating consists of two phases
(the inset at the top right corner of Fig.1a), generally identified as the β-phase (grey contrast) and the γ-
phase (light contrast).
The bulk NiCoCrAlY alloy with the same nominal composition as the as-sprayed NiCoCrAlY coating
was fabricated using spark plasma sintering (SPS). Briefly, the powder was packed in a column graphite
die with an internal diameter of 28 mm, and densified at 1050°C using a SPS system (FCT-HP D25/4-
SD) in vacuum (< 100 Pa ) at a load of 31 kN for 10 minutes. The as-sintered NiCoCrAlY alloy is fully
dense (Fig.1b) and consists of a β-phase (dark contrast) and a γ-phase (light contrast).
Rectangular (15 × 8 × 1.8 mm3) and disc (28 mm in diameter and 2 mm thick) samples were cut from
the bond-coated superalloy strips and the bulk NiCoCrAlY alloy bar, respectively, using a SiC abrasive
cutting blade in a precision cut-off machine (Accutom 5, Struers). The surfaces of the coatings and the
alloys were progressively ground and polished to a 1 μm finish to remove the existing surface asperities .
This step produced a similar initial root-mean-square roughness (~ 0.05 μm) on all samples. Some of
the coatings were mechanically thinned to several thicknesses in order to assess the effect of the
4
coating thickness on rumpling development. A micrometer was used to monitor the thickness
reduction during the thinning process and an optical microscope was used to double-check the final
thickness of the cross-section of the coating edge. Subsequent oxidation experiments showed that
coatings with thicknesses less than 50 μm tended to have TGO cracking and spallation after cyclic
oxidation for a period of time (e.g. 50 1-hour cycles at 1150°C) and, therefore, were not used in this
study. In addition, Vickers micro-hardness indentations were placed into the surfaces of several samples
as markers so that the rumpling evolution of the identical surface areas with increasing number of
thermal cycles could be tracked.
2.2 Thermal treatment
Cyclic oxidation was performed in laboratory air (1.01 × 105 Pa) between room temperature and
1150ºC in a CMTM rapid cycle furnace. Each cycle consisted of 10 minute ramping, 1 hour “hot time”
at 1150ºC and 10 minute fan-assisted air quenching. Some additional sets of cyclic oxidation
experiments were performed with a shorter “hot time” (10 minutes) at 1150ºC in each cycle, while
other cycling parameters remained unchanged. A few oxidation experiments were also carried out
isothermally at 1150ºC. The use of different oxidation regimes is to assess the effect of thermal history
on rumpling development. Apart from thermal exposure in air, several coating samples were cycled at
1150ºC in vacuum where the partial pressure of oxygen was significantly reduced (~ 1 Pa). This was
achieved by wrapping the samples in FeCrAlY foils (Goodfellow, UK) and then sealing them in quartz
tubes in vacuum. Due to its great resistance to thermal shock, the quartz tube was intact throughout
thermal exposure and therefore the vacuum was maintained.
2.3 Characterisation methods
The surface topography of the samples were characterised using a non-contact optical profilometer
(MicroXAM, KLA-Tencor). The profilometer provided digital images in the form of surface height, z,
as a function of lateral position, x and y. The rumpling magnitude is characterised by the root mean
square roughness, Rq, given by:
𝑅𝑞 = √1
𝑛∑ (𝑍𝑖 − �̅�)2
𝑛
𝑖=1 (1)
5
where n is the number of total data points, Zi is the height of each point and �̅� is the mean height for
the entire measured area. The images obtained were leveled to eliminate the effect of any macroscopic
slope on the roughness calculation. Since the rumpling patterns show substantial periodicity, the
wavelength, λ, of rumpling was characterised by the average distance between the adjacent undulation
peaks or valleys, which is given by:
𝜆 = 2𝜋
∆𝑅𝑞 (2)
where ∆ is the root mean square slope of the surface, given by:
∆ = √1
𝛿2𝑛∑ ⌈𝑍𝑖+1 − 𝑍𝑖⌉2𝑛𝑖=1 (3)
and 𝛿 is the spacing between the points (0.02 μm). At least 10 locations were recorded for each sample.
During thermal cycling, samples were removed from the rig at specific intervals for roughness
measurements and then returned for further cycling.
Following the surface roughness and stress measurements, the surface of the samples was carbon-
coated and examined using a scanning electron microscope (SEM, Quanta 650, FEI). SEM images of
the surface rumpling morphology were captured at the same tilt angle (60°) so the extent of rumpling
could be directly compared. Finally, the samples were cross-sectioned using a SiC abrasive cutting blade
and prepared in cross-section following conventional metallographic procedures. The microstructure
of the cross-sections was then examined using a combination of SEM and electron backscatter
diffraction (EBSD, NordlysNano, Oxford Instruments).
3 Results
3.1 Effect of thermal history and coating thickness on rumpling
Fig.2a-c shows the surface morphologies of three NiCoCrAlY coatings after 50 1-hour cycles. No
TGO spallation is observed throughout the oxidation. The initially planar coating surfaces rumple after
50 1-hour cycles, but the extent of rumpling shows a dependence on the thickness of the coating: the
rumpling becomes more pronounced as the thickness of the coating decreases from ~ 180 to ~ 60 μm.
The observation is further illustrated by the profilometer images in Fig.3 in which the surface
6
topographies (Fig.3a-b) and typical height profiles of the line segments along the lateral direction
(Fig.3d) , together with the corresponding Rq values, of a thin (~ 60 μm) and a thick (~ 180 μm)
coating after 50 1-hour cycles are shown. The roughness of the thin coating is about twice that of the
thick coating, yet the wavelengths (~ 60 - 70 μm, Table 2) are more or less the same. For comparison,
the coating surfaces, irrespective of the coating thickness, show far less roughening after 50 hour
isothermal oxidation (e.g. Fig.2d). This indicates that, for the same cumulative exposure time at 1150°C
and at least for the oxidation time used in this study, the bond coat is more susceptible to rumpling
under cyclic oxidation. The statistical rumpling amplitudes and wavelengths, associated with their
standard deviations, of the samples under different experimental conditions are listed in Table 2.
Fig. 4a shows the surface morphology of a coating (~ 60 μm thick) after 50 10-minute cycles. While the
coating surface rumples after 50 10-minute cycles, the extent of rumpling is less pronounced compared
with that after 50 1-hour cycles (Fig.2c). This observation could be further illustrated by the
profilometer images in Fig.3 where the surface roughness (Fig.3a and c) and profiles of line
interceptions (Fig.3d) after 50 1-hour and 10-minute cycles are compared. In short, the observation
indicates that for a given cycle number the exposure time at peak temperature plays an important role
in rumpling growth. This suggests that for a given cycle number and coating thickness, a larger net
TGO growth promotes rumpling development.
3.2 Rumpling of bulk NiCoCrAlY alloys
In contrast to the coating deposited on the superalloy substrate, the surface of the bulk NiCoCrAlY
alloy shows less tendency to rumple after cyclic oxidation. This is further illustrated in Fig.4b in which
the surface morphology of the alloy after 50 10-minute cycles is presented. The roughness, Rq, on the
coating surface (Fig.4a and Table 2) is about 2.5 times that on the alloy surface (Fig.4b and Table 2).
This indicates that while cyclic oxidation can produce rumpling on a bulk NiCoCrAlY as well, the
rumpling magnitude is considerably smaller than that of the coating subject to the same thermal history,
which suggests that the presence of the superalloy substrate is essential to rumpling development.
Observations on rumpling of the bulk NiCoCrAlY alloys after 50 1-hour cycles are not shown since
the samples have been through catastrophic TGO spallation and reoxidation, thus impeding the
7
application of the profilometer for the quantitative assessment of roughness. A possible reason for this
phenomenon is that the TGO cannot undergo enough stress relaxation as surface rumpling of the bulk
NiCoCrAlY alloys is small, the TGO stresses alleviate through spalling at a sufficiently high level of
thickness [7]. Another possible reason could be embrittlement of the TGO/alloy interface and
subsequent low interfacial strength caused by impurity segregation (e.g. Sulphur) during oxidation.
However, future studies of interface properties are needed to confirm this.
3.3 Evolution of rumpling with cyclic oxidation
Fig.5a-e presents a series of profilometer images recorded from the identical regions of a coating
surface (~ 60 μm thick) to illustrate the evolution of rumpling with increasing number of 10-minute
cycles. The initially flat surface gradually roughens with increasing number of thermal cycles.
Comparison of these images reveals that once rumpling is initiated, the areas above the average surface
plane continue to bow up while the areas below the average surface plane continue to depress during
the subsequent cyclic oxidation. Fig.5f shows the evolution of the height profile of a line segment (the
white line in Fig.5a) across the area probed during cyclic oxidation. The height profiles clearly show that
the coating surface progressively roughens with cycling oxidation (Rq, calculated based on the
cross-sectional height profile using Eq (1), increases from 0.21 μm after 5 cycles to 0.87 μm after 50
cycles).
Compared with the coating samples, rumpling increment on bulk NiCoCrAlY alloy samples is less
significant with increasing cycle number, as shown in Fig.6a-d where profilometer images recorded
from identical regions are presented. The surface topography after 10 10-minute cycles shows a
relatively small change compared with that after 50 10-minute cycles. This is further confirmed by the
evolution of the height profile of a line segment (the white line in Fig.6a) with cyclic oxidation, as
shown in Fig.6e where the roughness slightly increases from 0.17 μm after 10 cycles to 0.30 μm after 50
cycles.
3.4 Microstructure of coatings
To date, numerous experiments and simulations conducted to study rumpling of bond coats have
shown that the thermomechanical properties (e.g. coefficient of thermal expansion (CTE), creep and
8
yield strength) of the coatings greatly affect rumpling. These properties, in principle, are dictated by the
intrinsic microstructure of the coating. This relationship between the microstructure and properties
becomes even more complicated for a bond coat deposited on a superalloy substrate as the
microstructure of the coating constantly changes with thermal exposure. Therefore, characterisation of
the microstructure of the NiCoCrAlY coating and how it evolves with thermal cycling are of
importance to understanding its rumpling development.
Fig.7a-b shows the cross-sectional microstructure of a thick (~ 180 μm) and a thin coating (~ 60 μm)
after 50 1-hour cycles. TGO (predominantly α-Al2O3) scales with a thickness of ~ 3.5 μm form on
both coatings. The thickness of the TGO scale is fairly uniform over the coating surface, suggesting
that rumpling could not be induced by oxide intrusions or variation of the oxide growth rate from
place to place. Furthermore, the uniform TGO thickness over the bond coat surface suggests that the
top surface of the TGO accurately follows the undulations of the bond coat, which in turn confirms
that the Rq of the TGO is a good indicator of rumpling. Due to the selective oxidation of aluminium
and the interdiffusion between the coating and the substrate, the β-phase near the TGO/coating
interface and the coating/substrate interface decomposes to the γ-phase, which is also confirmed by the
EBSD phase-contrast maps in Fig.7c and d. As a matter of fact, for the thin coating the aluminium
depletion becomes so extensive that the β-phase could be hardly observed. Quantitative analyses (Fig.7e
and f) based on cross-sectional images using ImageJ software show that the changes of the volume
fractions of both phases follow a parabolic law, but the volume fraction of the β-phase in the thin
coatings (~ 60 μm) decreases around 2.5 times faster than that in the thick coatings (~ 180 μm), which
suggests that the thin coatings are less resistant to degradation during thermal cycling.
3.5 Rumpling of coatings after thermal cycling in vacuum
Fig.8 shows the surface and cross-sectional morphologies of a coating (~ 60 μm thick) after 50 1-hour
cycles in vacuum. The coating surface rumples (Fig.8a and d), yet the surface roughness (Rq ≈ 1.09 μm,
Fig.8a and c, Table 2) is lower than that after cycling in air (Rq ≈ 1.26 μm, Fig.2c and 3a, Table 2). The
thickness of the TGO scale is about 1.6 μm (Fig.8b), which is approximately half the thickness of the
TGO scale generated in air for the same exposure time (~ 3.5 μm). No β-phase is observed on the
9
cross-section of the coating after 50 1-hour cycles (Fig.8d). Fig.8e shows an inverse pole figure EBSD
map of the coating after 50 1-hour cycles where the crystallographic directions of the γ-grains parallel
to the RD direction (coating surface normal) are colour-coded. Only the γ-phase is indexed throughout
the entire coating. No correlation has been found between the concave/convex areas and the
crystallographic orientations of the metal grains underneath the TGO, which suggests that rumpling is
not associated with the crystallographic anisotropy of the γ-grains, whether through a dependence on
elastic modulus, yield strength, creep strength or diffusivity. These microstructural features (e.g. grain
size, grain morphology and phase composition) are extremely similar to that of the thin coating (with
the same thickness) after 50 1-hour cycles in air (Fig.7b and d). This suggests that, under the same
thermomechanical history (e.g. 50 1-hour cycles in air and vacuum), the difference in rumpling
magnitude between cycling in air and in vacuum is not associated with the microstructural difference of
the coatings. Therefore, the lower surface roughness after cycling in vacuum suggests that the TGO
growth is an important factor for rumpling development.
4 Discussion
4.1 Fundamental rumpling mechanism
In terms of the geometric profile, surface rumpling indicates overall lateral lengthening of the TGO
relative to that on a flat surface. The lengthening of the TGO could be associated with a couple of
mechanisms. One is the lateral growth of the TGO; the other is the out-of-plane distortion of the
coating surface, resulted from stress-driven plastic deformation of the coating. Unless the TGO
delaminates from the bond coat, TGO lengthening must occur along with plastic deformation of the
bond coat. Both mechanisms change the total strain energy of the system. The observations in this
investigation are used to evaluate the relative contribution of these mechanisms on the observed
rumpling behaviour.
The observations above have shown that, for a given cycle number and bond coat thickness, a longer
dwell time at the peak temperature produces a higher surface roughness and cycling in air produces a
higher surface roughness than that in vacuum. This indicates that the TGO growth is important to
rumpling development. However, the facts that isothermal oxidation produces limited surface
10
roughness on NiCoCrAlY coatings (Fig.2d) and the rumpling magnitude (Fig.4b) or increment rate
(Fig.6) of the NiCoCrAlY alloy during cyclic oxidation is relatively low during thermal cycling suggest
that the lateral TGO growth stress alone or thermal ratcheting of TGO (which involves both TGO
growth stress and TGO/alloy thermal mismatch stress) on an unsupported alloy alone is not sufficient
for accounting for the substantial rumpling development on the NiCoCrAlY coatings during cyclic
oxidation. This in turn indicates that the thermal mismatch between the coating and superalloy
substrate plays a crucial role in rumpling development (see Table 3 for CTEs of the NiCoCrAlY
coating and Hastelloy X substrate). However, the coating/substrate thermal mismatch alone could not
explain some observations in this study either. One implication of this mechanism is that, as long as the
exposure time at the peak temperature is sufficient for fully relaxation of the stress in the coating,
rumpling would only depend on the cycle number but not the dwell time at the peak temperature. On
the other hand, it is believed that full relaxation of the stress in the NiCoCrAlY coating should occur
within 10 minutes as previous stress measurements at 1150 °C have shown that the stress in this
coating is close to 0 after holding at this temperature for 10 minutes [28]. This is inconsistent with the
observations in this study, which show that the rumpling magnitude increases with increasing dwell
time at the peak temperature for a given cycle number. Furthermore, if the coating/substrate thermal
mismatch is the only single driving force for rumpling, thermal cycling in vacuum would be expected to
produce a larger, or, at least, an equivalent, rumpling magnitude than cycling in air (for the same coating
thickness, dwell time at peak temperature and cycle number). This is because the TGO scale in vacuum
is thinner than that in air and therefore has a lower bending stiffness, which makes it easier to deform
to accommodate the distortion of the coating surface. However, this is also contradictory to the
experimental observations in this study. Therefore, it is suggested that rumpling of the NiCoCrAlY
coating is driven by both TGO lateral growth and coating/substrate thermal mismatch.
4.2 Effect of coating thickness on rumpling development
The observations in this work also demonstrate that the rumpling magnitude of the NiCoCrAlY bond
coat increases with decreasing coating thickness. This dependence of the extent of rumpling on the
coating thickness is in agreement with the work of Deb et al. However, previous experimental [13] and
modeling work [29] has suggested that maximum rumpling occurs at intermediate bond coat thickness.
11
The reason for the discrepancy between these observations is explored below.
The thermal misfit stress ( ) generated in the bond coat upon cooling is given by (assume no creep
relaxation):
=𝐸 (𝛼 − 𝛼𝑆)∆𝑇
(1 − 𝑣 ) + (1 − 𝑣𝑆)ℎ 𝐸 /(ℎ𝑆𝐸𝑆) (4)
where E is the elastic modulus; α is the CTE; ΔT is the temperature drop; v is Poisson’s ratio and h is
the thickness. The subscripts “BC” and “S” refer to the bond coat and substrate, respectively. As the
thickness of the substrate is about 10 times or more thicker than that of the coating, the stress in the
coating is expected to be nearly independent of the coating thickness according to Eq.(4). For instance,
using the CTE and elastic modulus data given in Table 3 and volume fractions in Fig.7e and f (the
elastic modulus and CTE of the coating are functions of the volume fractions of the γ-phase and β-
phase), the calculated residual stress ranges from 629.2 to 654.8 MPa in the thick (~ 180 μm) coating
and from 659.7 to 683.6 MPa in the thin (~ 60 μm) coating, respectively, throughout 50 1-hours cycles.
The difference in bond coat stress due to thickness variation for a given cycle number is less than 5%,
which is believed to be not the main reason for the appreciable increase in rumpling as the thickness of
the coating decreases.
In the modeling conducted by Balint et al. [29], the thermomechanical properties of the coating are
taken to be independent of time and thickness over the entire thermal cycling process. However, in the
case of the NiCoCrAlY/Hastelloy X system used in this study, the coatings, especially the thin ones, are
subject to extensive aluminium depletion and a substantial β to γ phase transformation after 50 1-hour
cycles (Fig.7). This degradation could lead to considerable changes of the thermomechanical properties
of the coatings and compromise the prediction of the simulations using time and thickness
independent material properties. Specifically, the degradation of the coatings reduces their creep
strength at high temperature. Creep data from previous experimental tests and simulations have shown
that the creep strain rate of the polycrystalline β-NiAl is about two or more orders of magnitude
smaller than that of the polycrystalline γ-Ni [30-32]. According to a number of modelling results [6,
33-35], the decrease of the creep strength of the coating would promote rumpling growth. Indeed,
12
experimental studies on coatings with high creep strength (e.g. Pt-diffused γ/γ’ bond coats [36-40] and
γ/γ’ EQ (Equilibrium) bond coats [20]) have shown negligible rumpling even after extensive thermal
cycling at sufficiently high temperatures. The decrease of the creep resistance becomes more
pronounced as the thickness of the coating decreases since a thinner coating undergoes a more
extensive aluminium depletion and subsequent β to γ transformation during thermal cycling (Fig.7e and
f). The observations in this work suggests that the relatively rapid loss of the strength of the thin
coating, caused by the depletion of aluminium, makes it more susceptible to rumpling growth and,
therefore, play an important role in rumpling development.
4.3 Comparison of rumpling between the NiCoCrAlY coating and β-(Ni,Pt)Al coating
This work provides an opportunity to compare the rumpling behaviour of the NiCoCrAlY coating
with that of the β-(Ni,Pt)Al coating which has been extensively studied in the literature. The Rq of the
NiCoCrAlY coating after 50 1-hour cycles at 1150 °C ranges from 0.71 to 1.26 μm, increasing with
decreasing coating thickness. The β-(Ni,Pt)Al coating, however, has a Rq ~ 2.3 μm after exposure under
the same thermal history [41], which is 2 to 3 times that of the NiCoCrAlY coating. The characteristic
wavelengths of the rumpling patterns on the two types of coatings, however, are more or less the same
(~ 70 μm) [41]. To understand the reason for this difference between the rumpling magnitudes of the
two coatings, it is necessary to compare the respective thermophysical properties of the two coating
systems first.
Table 3 lists the CTE, phase composition, elastic modulus and TGO growth stress (TGO yield strength)
of the β-(Ni,Pt)Al/René N5 and NiCoCrAlY/Hastelloy X coating systems given in the literature [42-
47]. For the NiCoCrAlY coating, as it has been through a substantial β to γ phase transformation
during cyclic oxidation at 1150 °C, the CTE of the coating depends on its specific phase composition
[48]. The same argument also applies to the variation of the elastic modulus of the NiCoCrAlY coating.
The β-(Ni,Pt)Al bond coat deposited on the René N5 superalloy substrate, however, shows few phase
transformations from β to γ even after 100 1-hour cycles at 1150 °C [12], and therefore the CTE of the
coating can be reasonably seen as a constant. On the other hand, previous stress relaxation tests have
shown that the β-(Ni,Pt)Al coating shows a slower creep rate than that of the NiCoCrAlY coating [49-
13
51]. This suggests that the β-(Ni,Pt)Al is more creep resistant than the two-phase NiCoCrAlY. The
TGO growth stress on the NiCoCrAlY coatings is calculated based on the combination of the room
temperature TGO stress and the thermal-elastic parameters of the TGO and Hastelloy X substrate [52].
As the magnitude of the TGO growth stress is a result of a dynamic competition between the stress
generated by the lateral growth strain and any concurrent stress relaxation processes, the TGO growth
stress is also a good approximation for the yield strength of the TGO at the growing temperature.
The comparison of the thermophysical properties in Table 3 shows that compared with the β-(Ni,Pt)Al
coating deposited on the René N5 superalloy substrate, the NiCoCrAlY coating deposited on the
Hastelloy X substrate shows a higher CTE mismatch with the substrate, a higher elastic modulus, a
lower creep strength and a higher TGO yield strength when thermally cycled between ambient and
1150 °C. All these features would lead to a higher rumpling magnitude of the NiCoCrAlY coating for a
given thermal history based on the predictions of the simulations reported in the literature [17, 33, 34].
The experimental results in this work, however, show an opposite trend with the prediction of the
simulations. The discrepancy, therefore, should be accounted for by some other factors.
In searching for the reasons for the difference in the rumpling magnitudes between the NiCoCrAlY
coating and β-(Ni,Pt)Al coating, it is helpful to recall other key parameters essential to rumpling
development. In most simulations conducted in the literature, the lateral growth strain of the TGO is a
key parameter governing rumpling growth and rumpling is shown to increase with increasing the strain
incorporated in the models [6].On the other hand, previous experimental studies have shown that the
addition of minor reactive elements (e.g. yttrium and hafnium) into FeCrAl alloys or β-(Ni,Pt)Al bond
coats largely reduces surface undulation of the growing oxide [15, 53]. This effect has been ascribed to
a smaller lateral growth strain in the TGO due to the segregation of reactive elements in the growing
TGO grain boundaries [53-57]. Hence, a likely explanation of the smaller rumpling magnitude of the
NiCoCrAlY coating is a smaller lateral growth strain of the TGO resulting from segregation of yttrium
at the growing TGO grain boundaries.
In order to examine if the argument above is true or not, it would be helpful to see if there is
segregation of yttrium at the TGO grain boundaries. To achieve this, the chemistry of the TGO grain
14
boundaries was analysed using energy dispersive spectroscopy (EDS, X-MaxN 80T, Oxford Instruments)
equipped on a transmission electron microscope (TEM, TecnaiTM G2, FEI). The TEM sample was
prepared by a focused ion beam (FIB, Quanta 3D, FEI) cutting through the TGO surface followed by
an in-situ lift-out technique. Fig.9 shows a bright-field image (Fig.9a) and the corresponding STEM
annular dark-field image (Fig.9b) of the TGO on a NiCoCrAlY coating (~ 180 μm thick) after 2 1-hour
cycles. The TGO is about 0.8 μm thick, with its microstructure featured by the outer portion composed
of equiaxed grains and the inner portion composed of columnar grains. An EDS linescan profile (the
inset in Fig.9b) across a couple of TGO grain boundaries shows that the signals from yttrium bump up
when the linescan intersects with either of the grain boundaries, suggesting the presence of yttrium
segregation at the TGO grain boundaries.
It should be finally noted that Jackson et al [58, 59], however, recently have shown that rumpling of Hf-
doped B2 bond coats is comparable to the Hf-free β-(Ni,Pt)Al coating during thermal cycling. The
reasons for this are not completely understood, but could be probably explained from two aspects. First,
apart from the thermomechanical properties of the coating system, rumpling is also affected by other
parameters (e.g. initial roughness and wavelength [6, 11, 17]). On the other hand, the sustenance of the
dynamic segregation of reactive elements at the TGO grain boundaries and its subsequent effect on
rumpling is dictated by the reservoir of the elements in the underlying coating and its depletion during
thermal exposure [60]. After exposed at 1150 °C or above for sufficiently long time, the reactive
elements in the bond coat could be depleted to an extent that its concentration is so low that its
beneficial effect on rumpling becomes insignificant during subsequent thermal cycling.
5 Summary
The surface of a NiCoCrAlY coating deposited on a Ni-based superalloy progressively roughens during
cyclic oxidation. The rumpling magnitude depends on thermal history, coating thickness and oxidation
atmosphere. Compared with the coating, the bulk NiCoCrAlY alloy with the same composition shows
a smaller tendency to rumple after thermal cycling. The coatings, especially the thin ones, experience
substantial degradation (e.g. β to γ phase transformation) induced by oxidation and coating/substrate
interdiffusion. The observations together suggest that rumpling of the NiCoCrAlY coating is driven by
15
a combination of the lateral growth of the thermally grown oxide and coating/substrate thermal
mismatch. It is suggested that the chemical degradation of the NiCoCrAlY coatings, associated with
their microstructural/mechanical degradation, promotes rumpling growth. This effect is more
pronounced in a thinner coating as it is through a more substantial degradation for a given thermal
history. Compared with the β-(Ni,Pt)Al coating deposited on the René N5 superalloy substrate, the
NiCoCrAlY deposited on the Hastelloy X superalloy shows a rumpling magnitude 2-3 times lower,
which is due to the smaller TGO growth strain resulted from segregation of yttrium at the growing
TGO grain boundaries.
Acknowledgment
The authors would like to acknowledge Nicholas Curry, Nicolaie Markocsan, Per Nylen from
Production Technology Centre, University West, Sweden for supply of the coating samples. The
authors are also grateful to Mr. David Gordon for the help in vacuum sealing and Professor Brain
Ralph for proof-reading the manuscript.
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18
Fig.1
Fig.1 Microstructure of the NiCoCrAlY bond coat and bulk NiCoCrAlY alloy: (a) optical image of the
cross-section of the as-deposited bond coat; the inset at the top right corner of Fig.1a is a high-
magnification backscattered electron (BSE) image showing that the bond coat consists of β-phase (grey
contrast), γ-phase (light contrast) and interfacial pores (dark contrast) between the NiCoCrAlY particles;
(b) BSE image of the as-sintered bulk NiCoCrAlY alloy; the dark contrast is the β-phase and the light
contrast γ-phase
19
Fig.2
Fig.2 Surface morphology of (a) a thick coating (~ 180 μm), (b) a coating with an intermediate
thickness (~ 120 μm) and (c) a thin coating (~ 60 μm) after 50 1-hour cycles; (d) surface morphology
of a coating (~ 60 μm thick) after 50 hours isothermal oxidation
20
Fig.3
Fig.3 Profilometer images of three coatings after cyclic oxidation: (a) a thin coating (~ 60 μm) after 50
1-hour cycles; (b) a thick coating (~ 180 μm) after 50 1-hour cycles and (c) a thin coating (~ 60 μm)
after 50 10-minute cycles. (d) height profiles of the white lines in a-c.
21
Fig.4
Fig.4 Surface morphology of (a) a NiCoCrAlY coating (~ 60 μm thick) and (b) a bulk NiCoCrAlY alloy
after 50 10-minute cycles
22
Fig.5
Fig.5 Profilometer images of a NiCoCrAlY coating surface after (a) 5 10-minute cycles, (b) 10 10-
minute cycles, (c) 20 10-minute cycles, (d) 35 10-minute cycles and (e) 50 10-minute cycles recorded at
identical regions, as illustrated by the indentation marks. (f) The evolution of the surface profile and
corresponding Rq of the line interception shown in a.
23
Fig.6
Fig.6 Profilometer images of a bulk NiCoCrAlY alloy surface after (a) 10 10-minute cycles, (b) 20 10-
minute cycles, (c) 35 10-minute cycles and (d) 50 10-minute cycles and recorded at identical regions, as
illustrated by the indentation mark. (e) The evolution of the surface profile and corresponding Rq of
the line interception shown in d.
24
Fig.7
Fig.7 BSE images of the cross sections of (a) a thick coating (~ 180 μm) and (b) a thin coating (~ 60
μm) after 50 1-hour cycles; (c-d) EBSD phase-contrast maps of the thick and the thin coating after 50
1-hour cycles, respectively. The γ-phase is coded with red and the β-phase blue; (e-f) evolution of the
volume fractions of the γ-phase and the β-phase in the thick (~ 180 μm) and thin (~ 60 μm) coatings,
respectively, with thermal cycling. The volume fractions are estimated from the area ratios of the two
phases by processing cross-sectional images of the coatings using ImageJ software
25
Fig.8
Fig.8 Surface and cross-sectional morphology of a coating (~ 60 μm thick) after 50 1-hour cycles in
vacuum: (a) surface morphology, (b) fractured cross-section of the TGO at the edge of surface, (c)
profilometer image, (d) general view of the cross section and (e) an EBSD map of the coating where
the crystallographic directions of the γ-grains parallel to RD are colour-coded
26
Fig.9
Fig.9 TEM analyses of the TGO after 2 1-hour cycles. (a) bright-filed image, (b) STEM annular dark-
field image. The inset in (b) shows an EDS linescan profile across a couple of TGO grain boundaries.
The signal from yttrium bumps up when the linescan intersect with the grain boundaries.
27
Table 1 Chemical compositions (wt. %) of Hastelloy® X superalloy and NiCoCrAlY coating
Ni Cr Co Al Y Fe Mo W C Mn Si B
Hastelloy® X Bal 22 1.5 / / 18 9 0.6 0.1 1* 1* 0.008* NiCoCrAlY Bal 17 23 12.5 0.6 / / / / / / /
* Maximum
Table 2 Roughness (Rq ) and wavelength (λ) of coatings after thermal cycling
Cycle regime Sample configuration Atmosphere exposed Rq (μm) λ (μm)
50 × 1-hour Coating (~ 180 μm thick) Air 0.71 ± 0.07 64.39 ± 3.52 50 × 1-hour Coating (~ 120 μm thick) Air 0.99 ± 0.09 66.46 ± 3.78 50 × 1-hour Coating (~ 60 μm thick) Air 1.26 ± 0.11 65.13 ± 4.56 50 × 10-minute Coating (~ 60 μm thick) Air 0.85 ± 0.08 67.16 ± 3.89 50 × 10-minute Bulk alloy Air 0.35 ± 0.06 69.23 ± 5.21 50 × 1-hour Coating (~ 60 μm thick) Vacuum 1.09 ± 0.11 66.17 ± 4.16
Table 3 Thermophysical properties of the β-(Ni,Pt)Al/ René N5 and NiCoCrAlY/Hastelloy X coating systems. α is the mean CTE over the temperature range between room temperature and 1150 °C. E is
the elastic modulus at room temperature. σ𝑇𝐺𝑂𝐺 is the TGO growth stress at 1150 °C, which is also
taken as the TGO yield strength at 1150 °C. The mean CTE of the Hastelloy X is obtained by fitting the mean CTEs given in Ref. [42] using a quadratic polynomial and extrapolating it to 1150 °C.
α ( ×10-6/°C) Phase
composition E (GPa)
𝛔𝑻𝑮𝑶𝑮 /TGO yield strength
(GPa)
NiCoCrAlY 19.3 ~ 20.3
[48] γ+β
155 ~ 200 [51]
~ 1 [52]
β-(Ni,Pt)Al 16.0 [48] β 117 [49] ~ 0 - 0.4 [43-46] Hastelloy
X 17.1 [42] γ 196 [42] /
René N5 17.2 [48] γ/γ’ 145 [47] /