Surface & Coatings Technology 207 (2012) 413–420

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Surface & Coatings Technology

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Morphological evolution of thermal barrier coatings with equilibrium (EQ) andNiCoCrAlY bond coats during thermal cycling

J.J. Liang, K. Matsumoto, K. Kawagishi ⁎, H. HaradaHigh Temperature Materials Unit, National Institute for Materials Science, Sengen 1-2-1, Tsukuba 305‐0047, Japan

⁎ Corresponding author.E-mail address: KAWAGISHI.Kyoko@nims.go.jp (K. K

0257-8972/$ – see front matter © 2012 Elsevier B.V. Alldoi:10.1016/j.surfcoat.2012.07.036

a b s t r a c t

a r t i c l e i n f o

Article history:Received 16 February 2012Accepted in revised form 7 July 2012Available online 16 July 2012

Keywords:Thermal barrier coating (TBC)Failure mechanismEQ coatingThermal cycling

Morphology features and failure mechanisms of thermal barrier coating (TBC) systems under thermal cyclingare very complicated as it is influenced by many factors. It is known that the imperfections existing along theinterface between top coat and bond coat play a significant part in determining TBC spallation features. In thepresent study, the bond coat surface was polished prior to top coat deposition to remove the imperfections. Inthis case, different failure characteristics occurred. In order to see whether bond coat compositions influencefailure mechanisms on the polished surface, two TBC systems were investigated. One is with a newly devel-oped EQ (equilibrium) bond coat, designated as TMBC-1; the other is a commercially available NiCoCrAlYbond coat.

© 2012 Elsevier B.V. All rights reserved.

1. Introduction

In order to improve the reliability and efficiency of critical airfoilcomponents in gas turbine engines, high-temperature protective coat-ings are required. The success of the coating is assessed by its abilityto remain in place, resist oxidation, avoid spallation and limit interdiffu-sion between the coating and substrate [1]. This is particularly impor-tant for the thermal barrier coating (TBC) system which is the mostadvanced and structurally-complicated coating system until now.

TBCs generally consist of a Y2O3-partially-stabilized ZrO2 ceramictop coat, a thermally grown oxide (TGO), a metallic bond coat, and asuperalloy substrate [2,3]. It is reported that the surface condition ofthe bond coat prior to top coat deposition plays a significant role in de-termining TBC spallation resistance. For electron-beam physical vapordeposited (EB-PVD) TBCs, in spite of with Pt-modified NiAl bond coats[4–6] or MCrAlY bond coats [1,2], during thermal cycling the TGO usu-ally undergoes displacement instabilities with the aid of interface im-perfections, e.g. ridges, and thus the interfaces of the top coat/TGO andthe TGO/bond coat become successively rougher (rumpling) withthermal cycling. The rumpling of the TGO is known as a majorfailure-inducing factor because it can promote local separation and de-lamination in the TGO and its vicinity, eventually leading to bucklingand spallation of the TBCs [7–9]. There is growing evidence that rum-pling is driven primarily by a combination of thermal expansionmismatch and creep relaxation in the coating system. Additionally,phase transformations at high temperatures, e.g. β-NiAl→γ-Ni [4,10],and concentrations of some minor elements (such as C and Hf) [11]also play a role in determining the TGO rumpling. According to these

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facts, it is concluded that the initial rough interface between the topcoat and the bond coat can lead to TBC final spallation bymeans of rum-pling during thermal cycling. So in order to improve the TBC perfor-mance, the surface roughness of the Pt-modified NiAl coatings shouldbe minimized [12]. It is reported that the surface roughness of MCrAlYalloys affects the microstructure and adhesion of the alumina scale[13]. But there is little information available on the spallation mecha-nism of the TBCs with polished MCrAlY bond coat, so the presentstudy will investigate the failure characteristic of a TBC system withpolished MCrAlY bond coat.

In the present study, an innovative bond coat material will also beapplied besides MCrAlY. It is termed EQ (equilibrium) bond coat anddeveloped by the National Institute for Materials Science (NIMS),Japan [14–16]. Fig. 1 shows a pseudo-binary phase diagram of the(Ni,X)–(Al,Y) system [16]. Candidate compositions of the EQ coatingsfor Ni-base superalloys can be selected along the γ–γ' tie-line of thebase alloy substrate. Points A and B are in totally thermal equilibriumwith the substrate alloy (S point); hence, the composition of the EQcoating can be selected, such as the Al-rich point B. Since the chemicalpotentials of the alloying elements in the EQ coating and the substratein thermodynamic equilibrium state are equal, the interdiffusion willbe reduced, and the secondary reaction zone (SRZ) formation betweenthe bond coat and the substrate will be suppressed. Generally, the EQcoating in the totally thermodynamic equilibrium state does not pos-sess adequate oxidation resistance, so its composition is usually mod-ified a little, such as the concentrations of Hf, Y and Re. Based onRené-N5 superalloy, an EQ coating, named TMBC-1, is developed. Pre-vious study [16] showed that TMBC-1 can be applied to several gener-ations of Ni-base single crystal superalloy substrate in consideration ofsuppressing SRZ formation. This is because the difference of the chem-ical potential of Al in the coatings and substrates, which contributes to

Concentration

Tem

pera

ture

Fig. 1. Pseudo-binary phase diagram for the (Ni,X)–(Al,Y) system.

414 J.J. Liang et al. / Surface & Coatings Technology 207 (2012) 413–420

SRZ formation, is much smaller for the TMBC-1 coating than the con-ventional coating. Therefore, besides the above mentioned MCrAlYbond coat, the performance and failure characteristic of the surfacepolished TMBC-1 bond coat on widely-used CMSX-4 superalloy sub-strate will also be investigated in the present study so that the feasibil-ity of TMBC-1 on CMSX-4 substrate can be tested and the experienceon the EQ coating designing can be accumulated.

2. Experimental procedure

Ni-base single-crystal superalloy, CMSX-4, was chosen as the sub-strate in the study. The substrate samples were in the form of disk-shaped coupons with ≤001≥ orientation in thickness direction (10 mmin diameter and 5 mm thick). All the superalloy samples were grit-blasted, and then two ~120 μm thick bond coats, TMBC-1(an equilibriumcoating) andNiCoCrAlYwere deposited by the lowpressure plasma spray(LPPS) technique under the same processing parameters. The composi-tions of the substrate and bond coats are listed in Table 1. Prior to TBC ap-plication, the bond coated samples were polished to #800 SiC paper toreduce the existing surface roughness. The ~150 μm thick top coat of8 wt.% Y2O3/92 wt.% ZrO2 (YSZ) was applied by electron-beam physicalvapor deposition (EB-PVD) at 1000 °C in an oxygen atmosphere with apartial pressure of 0.2 Pa. Due to the high temperature and oxygen atmo-sphere during ceramic deposition process, a thin pre-oxidized TGO layerof about 0.2–0.3 μm formed on the bond coats.

All the TBC sampleswere tested in their as-coated conditions. Thermalcycling tests were conducted at 1135 °C in a rig built for this purpose.Each thermal cycle consisted of a 1-hour air exposure at 1135 °C and a1-hour cooling to room temperature (about 22 °C) by fan-forced labora-tory air. The coupons were transported automatically in and out of thefurnace, and the heating and cooling rates were estimated to be about300 °C/min and 500 °C/min, respectively. Periodically samples werechecked visually during cooling in order to examine the surfacemorphol-ogy. After prescribed cycles, some samples were taken out of the rig andanalyzed using optical microscopy (OM), scanning electron microscopy(SEM) and electron probe X-ray microanalysis (EPMA). Since thermallycycled TBC specimens are very sensitive to post-testing handling, all thesamples prepared for OM, SEM and EPMA were immersed in stabilizingresine, and then cutting, grinding and polishing were conducted. Atleast three samples of each coating were tested to failure.

Table 1Nominal compositions (wt.%) of the superalloy substrate and the bond coats.

Co Cr Al Mo W Ta Re Ti Hf Y Ni

CMSX-4 9.0 6.5 5.6 0.6 6.0 6.5 3.0 1.0 0.1 – BalanceTMBC-1 6.2 4.0 8.1 1.0 4.5 9.9 – 0.4 0.1 BalanceNiCoCrAlY 21.5 17.0 12.4 – – – – – 0.7 Balance

3. Results

3.1. As-coated microstructure

The cross-sectional microstructures of the as-coated TBCs arepresented in Fig. 2. In both cases, the TBC samples comprise a~150 μm thick YSZ top coat and a ~120 μm thick bond coat. Thetop coat exhibits typical columnar structure with a thin layer of ran-domly oriented, equiaxed grains at the bottom of the top coat. Beneaththe top coat, a thin pre-oxidized TGO layer (about 0.2 μm) is observed,which shows itself as a black line. In the two bond coats indicated inFig. 2, some pores exist. These pores originate from the deposition pro-cess. Discernible differences can be seen in the phase constituents of thebond coats fromFig. 2. The TMBC-1 bond coat is composed of one phase,which is designed to be γ'-Ni3Al, while there are two phases present inthe NiCoCrAlY bond coat, Al-rich β-NiAl and Al-poor γ-Ni, taking ondark and light appearances, respectively.

3.2. Thermal cycling life

The average cycles to failure (defined as approximately 50% spallationof the top coat by area) for the TBCs are presented in Fig. 3. Fig. 3 also in-cludes the spallation lives of TBCs on the CMSX-4 superalloy substratewith other types of bond coat materials, Pt-diffusion, low-temperaturePt–Al, and high-temperature Pt–Al, from Ref. [17]. It can be seen thatthe TBC spallation life was found to be slightly dependent on the type ofthe bond coat. TMBC-1, the firstly developed EQ coating, did not possessa superior spallation resistance when applied on the CMSX-4 substrate.

3.3. Post-failure surface microstructure

Although the two investigated TBC systems have similar thermalcycling lives, they exhibit distinct microstructural features on the

Fig. 2. The microstructures of the TBC systems in the as-coated condition: (a) TMBC-1bond coat, and (b) NiCoCrAlY bond coat.

0

50

100

150

200

250

300

350

400

450

500

NiCoCrAlYTMBC-1

Num

ber

of 1

135o C

Oxi

datio

n

Cyc

les

to F

ailu

re In the present study In Wu's study [17]

HT Pt-AlLT Pt-AlPt-difussion

Fig. 3. Spallation lives of the TBCs with different bond coats for 1-hour cyclic oxidationtesting at 1135 °C.

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post-failure surfaces, as shown in Fig. 4. The TMBC-1 bond-coatedsample exhibits a dark, clear spallation surface with little ‘scratches’dispersed on it. The ‘scratches’ are indicated by arrows in Fig. 4a andb. The spallation surface of the NiCoCrAlY bond-coated TBC, however,shows light color, and lots of dark ‘branches’ spreading on it (indicatedby arrows in Fig. 4c and d).

EPMA element distribution mapping (Fig. 5) illustrates that O, Al, andNi etc. cover the whole exposed spallation surface of the TMBC-1 bond-coated TBC, with alumina ‘scratches’ on it. This implies that the TGO re-mains, at least partially in-thickness, attached to the bond coat. The ele-ment distribution of the spallation surface for the NiCoCrAlY bond coatis different from the TMBC-1 (Fig. 5). The light region of the surface is cov-ered byNi, Co and Cr etc., and the ‘branch’ regionwith cracks at the centeris rich in Al andO, indicating that alumina ‘branches’ are embedded in thebond coat. It is reminded that all the bond coat surfaces prior to top coatdeposition were polished to remove imperfections and planarize it. Sothe formation of the oxide ‘scratch’ or ‘branch’ on the spallation surfacemust not be induced by the pre-existing imperfections.

Fig. 4. Post-failure macro-photographs of TBC surface after therm

3.4. Surface microstructure evolution during thermal cycling

In order to check the microstructural change and reveal the alumi-na ‘scratch’ and ‘branch’ formation mechanisms of the two investigat-ed TBC systems, half-coupon TBC samples from the same batch as thewhole ones were then tested under the same thermal cycling condi-tion. At this time, the top surface of the TBC-coated samples was checkedand recorded every day using optical microscopy (OM) during coolingdown. The results are presented in Figs. 5–7. It is seen that during thermalcycling, the top surface of the two TBC systems show different morpho-logical evolutions.

3.4.1. NiCoCrAlY bond coated TBCOn the top surface of the NiCoCrAlY sample some “line” regions with

slightly lighter contrast, noted as ‘white lines’ in the present paper, appearin a few tens of cycles, and they get a little longer andmuchwider duringfollowing cycles (Fig. 6). Fig. 7 shows the characteristic ‘white line’ athigher magnifications. One can see that cracks extend along the centerof the ‘white line’ in the samples after 100 cycles (indicated by arrowsin Fig. 7). Also, on the top coat surface of the failed sample (249 cycles)there are some regions that display relatively white color, marked by cir-cles in Figs. 6d and 7d. These regions exhibit similar color as the ‘whiteline’ and lie next to the edge of spalled region, indicating that separationbetween the TBC layers probably have occurred in this kind of regions.

3.4.2. TMBC-1 bond coated TBCFrom a macro-scale, the surface morphology of the TMBC-1 sam-

ple does not show observable change until spallation happened (Fig.8a and c); but when observed at a relatively higher magnification, asmall amount of cracks could be distinguished propagating in the ce-ramic (Fig. 8b). From the post-failure structure shown in Fig. 8d, it isseen that the spalled region exhibits dark color and a few tiny alumi-na ‘scratches’ distributes on the dark surface.

3.5. Cross-section microstructure evolution during thermal cycling

Morphological features of the TGO play a key role in the TBC failure, sothe development of such features during thermal cycling was also

al cycling at 1135 °C. (a), (b) TMBC-1, (c), (d) NiCoCrAlY.

Fig. 5. EPMA mappings illustrating the main element distribution on the post-failure surface of the TBC systems.

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examined. Figs. 8–11 show the cross-section SEM images illustrating themorphological features of the TGOs after various thermal cycles. Basedon these SEM images, a few observations and deductions can be madewith regard to the thermal cycling induced interfacial degradation.

3.5.1. NiCoCrAlY bond coated TBCDuring early thermal cycles the TGO in the NiCoCrAlY bond

coated sample undergoes localized downward displacement insta-bility, manifesting as a local penetration of the TGO into the under-lying bond coat (Fig. 9). This kind of TGO penetration has beendiscovered in Pt-modified NiAl bond coats and referred to ‘ratcheting’ inMunn's study [10], but the following development mode of ‘ratcheting’is different from his results. The ‘ratcheting’ progresses on a cycle-by-cycle basis, and must bring in stress in the superposed top coat that con-straint its growth. When the ‘ratcheting’ develop to an extent, through-thickness vertical cracks emerge along the inter/intra-column of the topcoat. The formation of the through-thickness vertical cracks leads to ex-cessive oxidation of the bond coat right below it. Thus, in the successive

Fig. 6. Macro-photographs of the TBC surface with NiCoCrAlY bond coat after (

cycling, the adhesion of the TGO to the bond coat in the vicinity of the‘ratcheting’ is reduced, and separations along this interface are initiatedandpropagated.When thenearby separationmeets each other, spallationoccurs. It is worth noting that the excessive oxidation of the bond coatleads to a small hollow area (cavity) surrounded by oxide in theratcheting. Also, within the bond coat around the ratcheting, the num-ber of voids has increased, compared with that observed in the as-coated state or in an undamaged region (Fig. 9). Within the voids,oxide is observed.

Based on the microstructural analysis of both the top surface and thecross section, the damage accumulation process of the NiCoCrAlY samplecanbe recited. In the early thermal cycling the TGO ‘ratcheting’ comes intoformation and lead to local debonding around it; the local debondingshows itself as the thin ‘white line’ on the top surface. As thermal cyclingproceeds, the debonding regions become bigger and vertical cracks formin the top coat, which can be seen from the much wider ‘white lines’ andcracks that extending along the ‘white lines’. As the damage progresses,the debonding grows and links together until reaching a critical size for

a) 48 cycles, (b) 100 cycles, (c) 160 cycles, and (d) 249 cycles at 1135 °C.

Fig. 7. Optical micrographs of the TBC surface with NiCoCrAlY bond coat after (a) 48 cycles, (b) 100 cycles, (c) 160 cycles, and (d) 249 cycles at 1135 °C.

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spontaneous bucking and spalling, indicated by the relatively white re-gion in Figs. 6d and 7d. The top coat, with the TGO attached to it, spallsfromthebond coat, leaving the excessively-oxidized ‘ratcheting’,manifestas the ‘branch’ on the post-failure surface, embedded in the bond coat.

In addition, besides the ratcheting, the TGO/bond coat interface in theNiCoCrAlY sample exhibits significant non-planarity. Since the bond coatsurface was polished to 800# SiC paper, these thickness heterogeneitiesevidently developed during the cyclic oxidation. Closer examinationshows this kind of heterogeneities is associated with entrained Y2O3/YAG precipitates having columnar morphology (Fig. 10). It has been sup-posed [18] that the Y2O3/YAG acts as a preferred channel for rapid inwarddiffusion of oxygen, resulting in locally thicker regions of the TGO (indi-cated by arrows in Fig. 10). Sohn et al. reported that embedded oxide,originating from shot-peening or heat-treatment of MCrAlY bond coatsprior to top coat deposition, leaded to formation of Ni/Co rich oxide parti-cles in the TGO, and causedfinal TBC failure [2]. In the present study, sincepolishing was conducted before top coat deposition, no such Ni/Co richoxide particle is found. But composition analysis by EPMA shows that

Fig. 8. Macro-photographs of the TBC surface with TMBC-1 bond

there is a thin (Ni, Co, Cr and Al)-rich oxide layer in the TGO next to thetop coat; this layer is suspected to be composed of spinel. According tothe interfacial degradation process mentioned above, it is seemed thatthe formation of the spinel layer does not influence the spallation resis-tance of the TBC.

3.5.2. TMBC-1 bond coated TBCThe interfacial microstructure evaluation of the TMBC-1 sample ex-

hibits different features from the NiCoCrAlY sample (Fig. 11). It is seenfrom Fig. 11 that a uniform TGO layer is formed in the TMBC-1 sampleduring early thermal cycles (Fig. 11a), and after several tens of cyclesthe initially uniform TGO layer displays kind of thickness heterogene-ities in some regions (Fig. 11b). With thermal cycling, the thicknessheterogeneities (protrusions) get bigger and void formation takesplace in some of the protrusions (Fig. 11c). It is worth noting that theTGO protrusions in the TMBC-1 sample differ from the ‘ratcheting’ inthe NiCoCrAlY sample. SEM observation demonstrates that until100 cycles there is no through-thickness vertical crack in the ceramic

coat after (a), (b) 300 cycles, (c), (d) 335 cycles at 1135 °C.

Fig. 9. The microstructures of the TBC system with NiCoCrAlY bond coat after (a) 30 cycles, (b) 50 cycles, (c) 100 cycles, and (d) 299 cycles at 1135 °C.

418 J.J. Liang et al. / Surface & Coatings Technology 207 (2012) 413–420

top coat of the TMBC-1 sample. Additionally, there are some oxide par-ticles entrained in the bond coat just below the TGO; this indicates thatthe TGO may contain some defects that allow for O2 diffusion towardthe bond coat. The failure mode of the TMBC-1 sample is governedby the separation in the TGO, neither along the interface of the YSZand TGO nor the interface of the TGO and bond coat. This is illustratedby the fact that a very thin TGO layer is left attached to the bond coat asshown in the post-failure cross-section image (Fig. 11d). From Fig. 11dit can also seen that the TGO protrusion is left in the bond coat. Basedon the cross-section (Fig. 11) and top surface (Fig. 8) observations, it isspeculated that the alumina ‘scratches’ dispersing on the post-failuresurface should be associated with the TGO protrusions embedded inthe bond coat.

Additionally, when the TMBC-1 sample undergoes about 80% cycliclifetime, vertical cracks in the top coat begin to appear. In most cases,the cracks do not lead to excessive oxidation and local de-cohesionbelow it, as shown in Fig. 12. Analysis of both the surface and thecross-sectionmicrostructural features indicates that this kind of cracksmay not play a vital role in the TBC spallation.

Fig. 10. Cross section SEM image of the TBC NiCoCrAlY sample after 50 cycles, showingthe Y2O3/YAG precipitates.

4. Discussion

The microstructure features expatiated above illustrate that thepolished bond coat surface prior to deposition of the top coat doesnot cause observable TGO rumpling in both the TMBC-1 and theNiCoCrAlY samples upon 1-hour thermal cycling at 1135 °C. Spitsberget al. reported that initial non-planarities, for example, ridges on thePt-modified NiAl bond coat, act as imperfections that initiate displace-ment instabilities, and removal of these imperfections, by polishing,inhibits instability formation, changing the failure mechanism toedge delamination with cracking along the TGO/bond coat interface[19]. But in the present study although both the two polished TBC sys-tems also failed at or near the TGO/bond coat interface, they exhibiteddifferent spallation mechanisms from edge delamination. What ismore, the two TBCs showed different damage accumulation processesand failure features from each other.

In order to explain this case clearly, our findings should be placed inthe context of previous work in this field. As thermal cycling proceeds,the strain energy density in the TGO increases, which is associatedwith thermal expansion mismatch [20–23], TGO growth [24–26], andphase transformations in the bond coat [4,10]. The strain energy den-sity can lead to out-of-plain stress in the crest of the imperfections [27]and causes TGO rumpling. When the strain energy density that acts asdriving force for separations accumulates to a certain level, localizeddebonding develops at or near the imperfections. In the present studythe imperfections have been removed by polishing, so the TGO rumplingthat is facilitated by interfacial imperfections is inhibited. In theNiCoCrAlYsample, another stress relief mechanism, TGO ratcheting and thecorresponding break-through cracks in the ceramic, is discovered. In theTMBC-1 sample, TGO deformation, such as rumpling and ratcheting, issignificantly reduced, and thus the strain energy density in the TGO can-not be relieved via the TGO deformation. Since the TMBC-1 sample is un-able to experience stress relaxation by rumpling or ratcheting, stressalleviation happens through TGO spallation. It is known that the twoTBC systems in the study were identically prepared, so the differencesin the stress relievingprocesses and failuremodesmust arise from thedif-ference in the bond coat properties.

Fig. 11. The microstructures of the TBC system with TMBC-1 bond coat after (a) 30 cycles, (b) 50 cycles, (c) 100 cycles, and (d) 335 cycles at 1135 °C.

Fig. 12. Cross section SEM image of the TBC system with TMBC-1 bond coat after335 cycles, showing the vertical crack in the top coat.

419J.J. Liang et al. / Surface & Coatings Technology 207 (2012) 413–420

The principal deformation driving force is the strain energy density inthe TGO, in other words it is the residual compressive stress in the TGO. Itis reported that the residual stress comes from the thermal expansionmismatch, TGO growth, and phase transformations in the bond coat.The coefficients of thermal expansion (CTE) of the two bond coat mate-rials were determined bymeans of calculating the changes in their latticeparameters based on high temperature X-ray diffraction peak, and theCTEs of the TMBC-1 and NiCoCrAlY materials at 1135 °C are calculatedto be 15.7×10−6/K and 17.0×10−6/K, respectively. The CTE of aluminascale is about (10.0–10.1)×10−6/K [28]. So the thermal expansionmismatch between bond coat and TGO is much smaller in the TMBC-1sample. This means the residual stress in the TGO caused by thermal ex-pansion mismatch is reduced in the TMBC-1 sample. In addition, thechemical composition of the TMBC-1 is very close to that of the substrate,so Al depletion due to interdiffusion between coatings and substrate issuppressed; thus the phase transformation in the TMBC-1 bond coatcaused by Al depletion is significantly suppressed. So the phase transfor-mation induced stress is also decreased. Besides, the TGO thickness in theTMBC-1 sample is similar to that in the NiCoCrAlY sample, and thus theTGOgrowth stressmay not lead to significant difference in themagnitudeof the TGO residual stress. Therefore, the driving force for TGO deforma-tion ismuch smaller in the TMBC-1 sample than in theNiCoCrAlY sample.

What is more, the difference in the deformation resistance of the twobond coats also makes contribution to the distinct failure mechanism ofthe TBCs. The W, Mo, and Ta in the TMBC-1, which are added to get the“equilibrium” with the substrate, can increase the creep strength of theTMBC-1. It is reported thatHf doping can also substantially increase defor-mation resistance of the TGO and bond coats [11], thus the addition of Hfin TMBC-1 bond coat can also decrease the ‘ratcheting’ of the TGO. For theNiCoCrAlY bond coat, however, the thickness heterogeneities associatedwith Y-containing precipitates can act as interface imperfection and facil-itate the TGO ‘ratcheting’.

We also tested the spallation life of the TMBC-1 bond coat on TMS-138A superalloy under the same cyclic oxidation condition, and the re-sults showed that its cycling life was higher than 1200 cycles. The rea-son for the distinct thermal cycling lives of the TMBC-1 on CMSX-4 andTMS-138A substrates need further investigation. Since the spallationresistance of the TMBC-1 bond coat depend strongly on the substrate

composition, it is advisable to develop a new EQ bond coat forCMSX-4. To achieve this aim, several factors should be taken into con-sideration, such as the interdiffusion between the new bond coat andCMSX-4, the concentrations of the oxidation-resistance elements andreactive elements, and the high temperature deformation resistanceof the bond coat.

5. Conclusion

A comparative investigation of TBC failure mechanisms associatedwith thermal cycling at 1135 °C has been presented. Removal of the im-perfections that exist on the bond coat surface before top coat deposition,throughpolishing, inhibits rumpling formation and changes failuremech-anisms. The NiCoCrAlY sample was found to undergo significantratcheting during thermal cycling. Break-through thickness cracks inthe top coatwas observed just above the TGO ratcheting. No ratchetingwas found in the TMBC-1 bond coated TBC. This was largely due to thedifferent stress states in the bond coat that was caused by the differentproperties of the bond coatmaterials, e.g. coefficient of thermal expan-sion. It is necessary to develop unique EQ bond coatmaterial for specif-ic superalloy substrate in order to improve TBC spallation resistance.

420 J.J. Liang et al. / Surface & Coatings Technology 207 (2012) 413–420

References

[1] R.C. Pennefather, D.H. Boone, Surf. Coat. Technol. 76 (1995) 47.[2] Y.H. Sohn, J.H. Kim, E.H. Jordan, M. Gell, Surf. Coat. Technol. 146–147 (2001) 70.[3] Y.H. Sohn, K. Vaidyanathan,M. Ronski, E.H. Jordan,M. Gell, Surf. Coat. Technol. 146–147

(2001) 102.[4] V.K. Tolpygo, D.R. Clarke, Acta Mater. 48 (2000) 3283.[5] V.K. Tolpygo, D.R. Clarke, Acta Mater. 52 (2004) 615.[6] V.K. Tolpygo, D.R. Clarke, Surf. Coat. Technol. 203 (2009) 3278.[7] A.M. Karlsson, A.G. Evans, Acta Mater. 49 (2001) 1793.[8] A.M. Karlsson, C.G. Levi, A.G. Evans, Acta Mater. 50 (2002) 1263.[9] A.M. Karlsson, T. Xu, A.G. Evans, Acta Mater. 50 (2002) 1211.

[10] D.R. MUMM, A.G. Evans, I.T. Spitsberg, Acta Mater. 49 (2001) 2329.[11] V.K. Tolpygo, K.S. Murphy, D.R. Clarke, Acta Mater. 56 (2008) 489.[12] N.P. Padture, M. Gell, E.H. Jordan, Science 296 (2002) 280.[13] A. Gil, V. Shemet, R. Vassen, M. Subanovic, J. Toscano, D. Naumenko, L. Singheiser,

W.J. Quadakkers, Surf. Coat. Technol. 201 (2006) 3824.

[14] K. Kawagishi, A. Sato, K. Matsumoto, J. Ang, H. Harads, In: in: J. Lecomte-Beckers, M.Carton, F. Schubert, P.J. Ennis (Eds.), Proceedings of the 8th Liège Conference, 2006,p. 691.

[15] A. Sato, H. Harada, K. Kawagishi, Metall. Mater. Trans. 37A (2006) 789.[16] K. Kawagishi, A. Sato, H. Harada, JOM (2008) 31.[17] R.T. Wu, X. Wang, A. Atkinson, Acta Mater. 58 (2010) 5578.[18] D.R. Mumm, A.G. Evans, Acta Mater. 48 (2000) 1815.[19] I.T. Spitsberg, D.R. Mumm, A.G. Evans, Mater. Sci. Eng., A 394 (2005) 176.[20] M.Y. He, A.G. Evans, J.W. Hutchinson, Acta Mater. 48 (2000) 2593.[21] V.K. Tolpygo, D.R. Clarke, K.S. Murphy, Surf. Coat. Technol. 146–147 (2001) 124.[22] V.K. Tolpygo, D.R. Clarke, Surf. Coat. Technol. 200 (2005) 1276.[23] M. Bäker, J. Rösler, G. Heinze, Acta Mater. 53 (2005) 469.[24] R.J. Christensen, V.K. Tolpygo, D.R. Clarke, Acta Mater. 45 (1997) 1761.[25] V.K. Tolpygo, D.R. Clarke, Acta Mater. 46 (1998) 5153.[26] V.K. Tolpygo, D.R. Clarke, Acta Mater. 46 (1998) 5167.[27] H. Bhatnagar, S. Ghosh, M.E. Walter, Int. J. Solids Struct. 43 (2006) 4384.[28] J.A. Haynes, B.A. Pint, W.D. Porter, I.G. Wright, Mater. High Temp. 21 (2004) 87.