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A mechanistic understanding on rumpling of a NiCoCrAlYbond coat for thermal barrier coating applicationsDOI:10.1016/j.actamat.2017.02.003

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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

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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 (p.xiao@manchester.ac.uk).

mailto:p.xiao@manchester.ac.uk

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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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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] /