Surface & Coatings Technology 203 (2008) 449–455

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Surface & Coatings Technology

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Cyclic behavior of EB-PVD thermal barrier coating systems with modified bond coats

Uwe Schulz ⁎, Klaus Fritscher, Andrea Ebach-StahlDLR-German Aerospace Center, Institute of Materials Research, 51170 Cologne, Germany

⁎ Corresponding author. Tel.: +49 2203601 2543; fax:E-mail address: uwe.schulz@dlr.de (U. Schulz).

0257-8972/$ – see front matter © 2008 Elsevier B.V. Aldoi:10.1016/j.surfcoat.2008.08.056

a b s t r a c t

a r t i c l e i n f o

Available online 4 September 2008

Keywords:

The lifetime of thermal baoxide (TGO), and ceramic toby LPPS and EB-PVD (elect

Thermal barrier coatingThermally grown oxideEB-PVDBond coat

ron-beam physical vapour deposition) underneath conventional EB-PVD yttriastabilized zirconia top coats were investigated on three different substrate alloys. Several bond coattreatments such as polishing, annealing in vacuum, and grit blasting were employed in order to study effectson TBC life, and particularly the underlying mechanisms of TGO formation. Samples were thermally cycled at1100 °C and partly at 1150 °C. Spallation of the TBCs is mainly correlated with TGO formation that isinfluenced by bond coat type and pre-treatment. The longest lifetimes were achieved on a novel Hf-doped

rrier coating (TBC) systems depends on substrate, bond coat, thermally grownp coat. In the present paper NiPtAl bond coats as well as NiCoCrAlY(X) deposited

EB-PVD NiCoCrAlY-X bond coat owing to a differing TGO formation and failure mechanism. Activationenergies derived from lifetimes and test temperatures were calculated to identify key failure mechanismswithin these complex coating systems.

© 2008 Elsevier B.V. All rights reserved.

1. Introduction

Thermal barrier coatings (TBCs) offer the potential to significantlyimprove efficiencyof aero engines and stationary gas turbines for powergeneration. All their constituents that are bond coat, thermally grownoxide, ceramic top and even the substrate they protect are crucial forlifetime of the coatings [1,2]. State-of-the-art TBCs consist of a PtAldiffusion or an MCrAlY overlay bond coat and a ceramic top coatdeposited by electron-beam physical vapor deposition (EB-PVD) orplasma spraying. The EB-PVD process offers the advantage of a superiorstrain and thermal shock tolerance for the ceramic coatings due to theircolumnar microstructure. Although increasing gas temperatures led tothe development of alternative chemical compositions, the currentmaterial in use is still partially yttria stabilized zirconia (7YSZ).

Manufacture of TBC systems is a multi-step process where everydetail may influence TBC life. While TBC thickness and cycle length intesting seem to be of second order importance for life of MCrAlY basedTBCs [3], EB-PVD process parameters do influence TBC life. Therefore, inthe present study EB-PVD process parameters were kept in a narrowband to exclude an influence of processing on TBC life. Proper bond coattreatments prior toTBCdepositionoffer a largepotential to economicallyimprove TBC life [4–7]. During processing and in service, a thermallygrown oxide (TGO) layer forms as a result of bond coat (BC) oxidationwhich usually plays the most important role for the lifetime of the TBC.Failure in EB-PVD TBCs is almost always initiated at or near the TGO,mostly along the TGO/BC interface. Although literature on individualcoating types is numerous, direct comparison between them is scarce

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[8–11]. It is obvious that testing temperature has a great influence oncyclic TBC life, with a lifetime reduction of about a factor of 10 for every100 °C temperature increase under cyclic testing conditions [11].

The present paper compares the cyclic life of standard 7YSZ topcoats on a variety of differently processed and treated bond coats. Forselected versions the role of testing temperature was studied. Specialattention was paid to the influence of rare earth elements such as Yand Hf in the bond coats on cyclic TBC life.

2. Experimental

Cylindrical samples of three different Ni-base superalloys wereused in the present study. Major differences involve the Ti and Mocontent in IN100 while CMSX-4 and Rene 142 are additionally alloyedwith W, Ta, Re and contrasting contents of Hf (see Table 1). All threealloys are used in turbine blade applications with the highesttemperature capability of the CMSX-4, followed by Rene 142 andIN100. Samples of 6 to 10 mm diameter received the following bondcoats and treatments (a–d) prior to TBC deposition that areadditionally summarized in Table 2 (for details of sample preparationand treatments see references for each treatment):

(a) 100 to 120 µm thick EB-PVD Ni-22Co–20Cr–12Al–0.1 to 0.2Y(wt.%) bond coats, designated EB-MCrAlY, were densified byball peening and subsequently vacuum annealed at 1080 °C for4 h at a pressure lower than 5⁎10−5 mbar [2]. A selected versiongot an exceptional smooth substrate finishing prior to bondcoat deposition (4000 grit finish). Only for this version, aslightly changed testing procedure was employed that reducedthe arrest time in ”humid” laboratory air for visual inspectionsat room temperature [12].

Fig. 1. Cyclic lifetimes of 7YSZ TBCs at 1100 °C on various state-of-the-art bond coatsusing common industrial practice for bond coat treatment prior to TBC deposition.

Table 1Chemical composition of substrate materials, measured by X-ray fluorescence (wt.%,Nickel is balanced)

Alloy Co Cr Al Mo W Ta Ti Hf Re C B Other

IN 100 DS 14 9 5 2.3 5 0.18 0.015 0.05 Zr, 1 VCMSX-4 9 6.5 5.6 0.6 6 6.5 1 0.1 3Rene 142 12 6.8 6.1 1.5 5 6.4 1.5 2.8 0.12 0.015 0.02 Zr

Concentration ofminor additions of C, B, and others are taken from nominal compositionsof the manufacturer of the alloys.

450 U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

(b) A 75 to 95 µm thick novel EB-PVD bond coat was applied bytwo-source EB-PVD in order to obtain a composition of Ni–23Co–23Cr–10.5Al–0.3 to 0.5Si–0.1Y–0.5 to 1.0Hf (designatedEB-MCrAlY-X). The variation in hafnium content was caused byvariable distances along the cylindrical samples from the Hf-source. Samples were subjected to the same vacuum heattreatment as the NiCoCrAlY bond coat (a).

(c) Approx. 50 to 80 µm thick commercial PtAl diffusion bond coatswere supplied by two different vendors (1) and (2). Thestandard treatment consisted of lightly grit blasting with220 mesh alumina prior to TBC deposition. For an alternativeversion of PtAl(1) the annealing atmosphere at 1080 °C for 4 hwas replaced by a mixture of argon and hydrogen at atmo-spheric pressure (designated PtAl(1)ArH) [3,10,13].

(d) NiCoCrAlY of the nominal composition Ni–20Co–12Cr–13Al–0.7Y with traces of Hf/Si (wt.%) was deposited by VacuumPlasma Spraying (VPS-MCrAlY) at research center Juelich FZJ.The standard treatment consisted of the common diffusion heattreatment at 1080 °C for 4 h followed by 870 °C for 20 h, both invacuum. After smoothing the bond coat by an industrial source,grit blasting was applied prior to EB-PVD TBC deposition. Analternative version had a reversed order of processing steps:smoothing by hand polishing followed by vacuum annealingprior to EB-PVD TBC deposition [6,14]. One version on Rene 142substrates followed the sequence diffusion heat treatment+smoothing+vacuum annealing at 1080 °C for 4 h.

Table 2Summa

Coating

PtAl(1)PtAl(1)AEB-MCrEB-MCrVPS-MCVPS-MCAPS

Versions (a) to (d) were given finally a 150 to 200 µm thick 7YSZceramic top coat by reactive EB-PVD in rotationmode at 12 rpm. Inorder to avoid an influence of EB-PVD processing parameters oncyclic TBC life, preheating of the samples, substrate temperature,chamber pressure and amount of oxygen were kept constant.

(e) An air plasma-sprayed top coat (APS) was applied on theidentical VPS bond coat with the common preparation steps ofvacuum annealing as described above for d) followed by APSdeposition [6].

While IN100 and CMSX-4 were used as base materials for all theabove coatings, Rene 142 was used only for selected coating systems(versions (a), (c1), and (d) with a slightly changed composition). Allsamples were cyclically tested at a holding temperature of 1100 °Cusing 1 h-cycles (50 min heating, 10 min forced air cooling to roomtemperature). Unless otherwise mentioned, a minimum of threesamples per condition was tested. Selected samples were additionally

ry of bond coat deposition method, applied BC treatment, and 7YSZ top coat deposi

designation Bond coat deposition method

and PtAl(2) Electroplating Pt+Al diffusionrH Electroplating Pt+Al diffusionAlY EB-PVDAlY-X EB-PVD two-sourcerAlY VPSrAlY improved VPS

VPS

tested at 1150 °C under the same cyclic conditions. Failure of the TBCswas defined as TBC spallation of an area with one dimension greaterthan 5 mm. Representative samples were cross-sectioned afterdeposition and after failure. They were prepared by conventionalmetallographic preparation techniques for examination by scanningelectron microscopy and energy-dispersive X-ray spectroscopy.

3. Results and discussion

3.1. State-of-the-art bond coats

In Fig. 1 cyclic TBC life is displayed for standard bond coattreatments. For all versions TBC life was longer on CMSX-4 than onIN100 with the exception of the EB-PVD bond coat that had a reversedorder. A ranking of cyclic TBC life on IN100 can be given as follows: EB-bond coatNPtAl(2)NPtAl(1), VPS-MCrAlY, APS-system. The ranking onCMSX-4 is similar with the only change for the EB-PVD bond coat: PtAl(2)NPtAl(1)NEB-bond coat, VPS-MCrAlY, APS-system. Note that thefully plasma-sprayed TBCs behaved nearly similar to the short life EB-PVD top coats despite of the inherent lower strain compliance of thissystem.

Cyclic TBC life of optimized versions that arose from severaldevelopment projects are depicted in Fig. 2. Again, on the optimizedEB-NiCoCrAlY bond coat TBC life is longer on IN100 than on CMSX-4while for the PtAl system (1) it is the opposite. TBC life on the vacuumplasma-sprayed bond coat was nearly the same on both substrates.Note the dramatic increase in TBC life of the optimized versions by afactor of approx. 2 (EB-bond coat on both substrates and VPS-MCrAlYon CMSX-4) up to a factor of 3 and higher (PtAl(1), VPS-MCrAlY onIN100) in comparison to the standard treatments given in Fig. 1. Thechemically modified EB-MCrAlY-X system is discussed separately inSection 3.2.

Two samples of four selected substrate–coating systems weretested at 1150 °C, using the same equipment and time regime as for

tion method

BC treatment Top coat

Grit blasting EB-PVD4 h at 1080 °C in ArH at ambient pressure EB-PVDPeening+4 h at 1080 °C vacuum EB-PVDPeening+4 h at 1080 °C vacuum EB-PVDDiffusion annealing, smoothing, grit blasting EB-PVDSmoothing, 4 h at 1080 °C vacuum EB-PVDDiffusion annealing APS

Fig. 3. Cross section of TGO of version IN100+NiCoCrAlY (improved)+ 7YSZ TBCafter 472 cycles at 1150 °C. SEM micrograph overview (a) and detail of outer TGO part(b). Bright particles in this layer are Y-rich oxides.

Fig. 2. Cyclic lifetimes of 7YSZ TBCs at 1100 °C on improved and novel bond coats. Theimprovements consisted of: annealing in ArH for 4 h at 1080 °C prior to TBC deposition(PtAl), 4000 grit polished samples and reduction of water vapour during roomtemperature inspection intervals (EB-MCrAlY), smoothing and reversed processingsequence (VPS-MCrAlY) and novel chemistry (EB-MCrAlY-X).

451U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

the 1100 °C tests. EB-PVD TBCs failed after the following averagedcycle numbers. Numbers in parenthesis give the 1100 °C results forcomparison; variations in sample numbers are marked. Activationenergies of lifetimes are given on the right.

a) on IN100 substrates:– PtAl(2) standard grit blasted 180 cycles (611); 397 kJ mol−1

– EB-MCrAlY improved 472 cycles—only one sample—(1422);358 kJ mol−1

b) on CMSX-4 substrate:–modified EB-MCrAlY-X 1272 cycles—three samples were tested—(4693); 424 kJ mol−1

– PtAl(2) standard grit blasted 504 cycles (960); 209 kJ mol−1

Ranking and relative life of the coating systems are nearlyunchanged if testing temperature is raised by 50 °C. The absolutelife time is roughly one third at the higher temperature with theexception of PtAl on CMSX-4. Microstructure investigations revealedthat TGO formation, failure location and crack pattern are also similarat both testing temperatures. Therefore, only selected cross sectionsare shown here.

TGO formation after cyclic tests at 1150 °C is shown in Fig. 3 for theEB-NiCoCrAlY system. Significant visual differences between thestandard and improved versions EB-NiCoCrAlY did not exist. Asdetailed in [2,15], the TGO consists of two layers. The outer mixed zoneof less than 1 µm thickness is slightly porous and consists of alumina–zirconia; quite often in an off-plane arrangement of alternating areasof (i) zirconia-free alumina underneath yttrium-oxides that hadformed during vacuum annealing, and of (ii) areas containing thecommon mixed zone arrangement of zirconia particles embedded inalpha alumina. Because diffusion processes are accelerated at thehigher testing temperature, this off-plane arrangement was notclearly visible after 1150 °C testing. The thicker inner TGO zone isdense alpha alumina. The bright particles in this region (see Fig. 3) areY-rich oxides, commonly identified as yttrium-aluminates with amajority of the YAG phase. They grew slightly larger in size at thehigher testing temperature at the expense of their number. Spallingof the TBC occurred either between the bond coat and TGO oroccasionally between the dense alumina TGO andmixed zone. EB-PVDmanufactured NiCoCrAlY bond coats preserve a relatively flat TGOover the entire testing duration at both temperatures.

Due to space limitations, the PtAl and VPS-MCrAlY version are onlydescribed verbally here. The TGO on all versions NiPtAl showed signs

of rumpling with an increasing number of cycles. Previous investiga-tions have revealed that an initially flatter interface between the bondcoat and TBC delays TGO rumpling and, hence, increases TBC life [10].Martensite was observed in the PtAl bond coat occasionally afterprolonged testing. The platinum concentration dropped below 5 at.%in the coating region due to fast diffusion of this element along theconcentration gradient towards the substrate for all versions, howeverthis was accompanied by substantial thickening of the bond coatthrough from interdiffusion of elements from the substrate alloy. In allcases the TGO consisted of alpha alumina with rare indications for theformation of an outer porous mixed alumina–zirconia zone. TGOthickness at failurewas in the range between 5 and 7 µmwith no largedependence on testing time. This can be explained by the rapid initialoxidation, where a large portion of the TGO thickness is formed duringthe first 100 h of testing while subsequent TGO thickening of PtAlbond coats is extremely slow [3]. The fully plasma-sprayed system hadthe typical rough TBC bond coat interface and showed indication ofpartial white failure, i.e. cracks above the interface TGO to TBC withinthe TBC.

The two phase PtAl bond coats (1) and (2) differ in both, micro-structure and cyclic life. Although nominally identical, version (1) wasthicker and had a higher aluminum content while version (2) had ahigher Pt content close to the BC to TBC interface. The higher amountof platinum in version (2) after bond coat deposition seems toovercompensate the lower total coating thickness and the smallerreservoir of aluminum.

Fig. 4. Cross section of CMSX-4 with novel EB-PVD MCrAlY-X bond coat after TBCdeposition. a) overview, b) higher magnification of TGO. Bright particles are Hf-rich; if incontact to TGO it is oxidized and far away still metallic.

452 U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

A significant result of the current work is that the bond coatsperform differently on each superalloy. On the Ta free polycrystallinealloy IN 100 (see Table 1), the superiority of a clean bond coatcontaining metallic yttrium becomes clearly visible. Consequently, along TBC life was achieved on IN100 with the EB-PVD bond coat.Previous work [12] has shown that among substrate elementadditions Ta had a negative influence on TBC life. The other substrateelements that differ between the used polycrystalline IN100 and thesingle crystal CMSX-4 alloy such as C, B, Re and W did not change EB-PVD TBC life on model cast alloys. It remains still unclear why areversed performance order of the coatings on the two substratesbetween EB-PVD and all other bond coats exists. It can be onlyspeculated that diffusion in the EB-PVD bond coat is the fastest for allinvestigated bond coats due to the small grain size that favors grainboundary diffusion and clean grain boundaries because of the lowestoxygen content in the EB-PVD coating (especially if comparedwith theVPS version). Tantalum may hence diffuse faster from the substratetowards the bond coat/TGO interface, where it seems to degrade TGOadherence. For PtAl it is agreed that diffusion of substrate elementstowards the surface is slowed down which would cause lessdetrimental effects of Ta from the SX alloy in the TGO or at the TGO/bond coat interface. Hence other effects arising from the SX alloy maydominate, leading to a longer TBC life of PtAl on CMSX-4 than onIN100 substrates.

While for the EB-bond coat it was not clear which factor wasmost important for lifetime enhancement of the improved version(a smoother substrate or reduced water vapor at room temperatureduring testing), the bond coat treatments on PtAl and on VPS-MCrAlY did clearly increase TBC life. The underlying mechanismsare summarized shortly. For the VPS bond coat, a reversed sequenceof processing steps assured formation of larger amounts of yttria onbond coat grain boundaries during vacuum annealing, giving risefor formation of a strong off-plane mixed zone [6,16]. This optimizedTGO microstructure provides a longer TBC lifetime, most pro-bably due to a higher effectiveness of the active rare earth elementyttrium.

The influence of annealing PtAl in ArH atmosphere is discussed indetail for the CMSX-4 substrate elsewhere [13] and operates similarlyfor the IN100 substrate samples. The oxygen partial pressure duringthis annealing is high enough for partial oxidation that leads to anoxide film consisting mainly of alpha alumina prior to TBC deposition.The EB-PVD top coat obviously adheres better on this TGO comparedto the one that forms on grit blasted surfaces during EB-PVDprocessing, leading to prolonged TBC life. In [17] an oxidation pre-treatment has also prolonged TBC life on PtAl bond coats by a factor oftwo. Vacuum annealing of PtAl did not provide an adequateatmosphere in order to achieve full coverage of the bond coat surfaceby alumina [3,10,13].

3.2. Novel EB-bond coat

The best lifetime improvements were achieved on the CMSX-4substrate by doping standard EB-PVD MCrAlY compositions with Hf(EB-MCrAlY-X in Fig. 2). Although the measured aluminum content inthe as manufactured stage was slightly lower for the NiCoCrAlY-SiHfversion compared to the standard NiCoCrAlY (10.5 vs. 12 wt.%), smalladditions of the reactive element Hf did increase TBC life by a factor of15. A comparative test of this coating on a single IN100 sample,however, brought about only a marginal life time improvement whichis the topic of ongoing research.

In the as-coated state shown in Fig. 4 hafnium has already diffusedtowards the bond coat surface and is present on top or underneath theTGO. Bright particles are Hf-rich; if in contact to the TGO they areoxidized and away from the interface they are still metallic. In analogyto NiCoCrAlY it can be concluded that during vacuum annealing Hf ispartially oxidized. Due to the low concentration of Hf in the bond coat

(≤1 wt.% ), surface coverage with HfO2 is incomplete. The preferred Hfdiffusion path is along grain boundaries between the two phases ofthe bond coat that are based on beta-NiAl and gamma-Ni solidsolution, as can bee seen from the position of bright particles in Fig. 4.In analogy to the mechanism in yttrium dominated bond coats such asvacuum annealed EB-PVD [16] and VPS NiCoCrAlY [6], an off-planemixed zone has formed on EB-PVD NiCoCrAlY-SiHf. Underneath theupper bright hafnia particles pure alumina forms (Fig. 4). Adjacent tothese particles, a typical mixed zone consisting of zirconia particleswithin alumina is present. Its formation is assumed to follow the samereaction scheme as in NiCoCrAlY where solution of zirconia intransient alumina or in spinel type oxides occurs with subsequentprecipitation of zirconia when the TGO transforms into the stablealpha alumina polymorph [16,18,19].

TGO formation during cyclic testing and after TBC failure is for thenovel bond coat quite different from the Hf-free counterparts (Fig. 5).The major differences comprise (i) a broccoli-structured wavy inter-face between TGO and bond coat accompanied by a large and varyingTGO thickness of 12 to 30 µm, (ii) bright and mostly rounded hafniaparticles spread throughout the whole TGO thickness, (iii) occurrenceof short-distance cracks in the TGO visible parallel to the interface, and(iv) NiCoCrAl-spinel oxides spread in the TGO (medium gray areas).The TGO consists of an approximately 1 µm thick outer porous mixedzone and a thick inner dense TGO underneath.

The undulated TGO bond coat interface is caused by oxide pegs ofalumina that surround Hafnium oxide particles. Cracks appear already

Fig. 5. CMSX-4 substrate withmodified EB-PVDMCrAlY-X bond coat after a)–c) intermediate testing interval of 1584 cycles at 1100 °C and d) after TBC failure at 1490 cycles at 1150 °C.Bright particles in TGO are Hf-oxides, medium grey particles are NiCoCrAl-spinel oxides.

453U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

after intermediate testing times in cross sections. Although verycarefully prepared, it can be never conclusively declared that all crackswere already present prior to sample preparation. Cracks occurmainlyat the TGO bond coat interface where they bridge the TGO in areas ofoxide pegs that have a larger thickness, but also at various distancesfrom the TGO–TBC interface (Fig. 5b and c). The cracks cannot followthe wavy interface TGO/bond coat. It remained unclear what themechanism for crack arresting is, but all cracks are only a few 100 µmlong in maximum. Cracks cross hafnia particles or go aroundthem. From cross sections one would deduce that crack densityincreases within the TGO at the higher testing temperature, therebycreating a blocky TGO appearance. Although still speculative, thisspecific short-distance cracked TGO microstructure is believed to bebeneficial for the dramatically prolonged TBC life. One possibility isthat stresses within the TGO are lowered and therefore the crackdriving force stays below the critical one for TBC spallation overprolonged times.

Although quite often believed to be detrimental in a TGO, spineltype oxides embedded in alumina were not harmful in the presentinvestigation as they do not influence cracks or diffusion paths of e.g.oxygen. Neither silicon nor yttriumwas identified in the TGO region inany enriched phases, although they were both present in minorconcentration in the bond coat. Silicon was still present within thebond coat after testing while Y has most probably off-diffused into theTGO, but did not form larger oxide particles as common for NiCoCrAlYbond coats.

3.3. TBCs on Rene 142 substrates

On the Hf-containing superalloy Rene 142 selected bond coatswere applied and the following lifetimes at 1100° were recorded:

(i) 2327 cycles for VPS-MCrAlY of a slightly changed compositionNi–38Co–19Cr–10Al–0.4Y (measured values of Amdry 995,supplied by an industrial vendor) with grit blasting treatmentand 1138 cycles with vacuum heat treatment, both treatmentswere applied after the smoothing step;

(ii) 2000 cycles for standard EB-NiCoCrAlY (test has been stoppedafter this time without TBC failure);

(iii) 988 cycles for PtAl(1) using standard grit blasting prior to TBCdeposition.

TBC lifetimes are dramatically longer for both MCrAlY versionstested and about twice that for PtAl(1) on this substrate compared toIN100 and CMSX-4 substrates and identical bond coat treatments. Forthe VPS system results prove that a more conventional sequence ofprocessing steps (diffusion annealing, smoothening, followed byvacuum annealing) is less favorable compared to grit blasting sincesmoothing removes most of the yttria formed during vacuumdiffusion annealing [6]. It remains for further research to confirmthat a reversed sequence of processing steps as shown for CMSX-4 andIN100 (Fig. 2) would give a longer life of VPS bond coats on Rene 142 aswell.

Fig. 6. SEM cross section of Rene 142+VPS-MCrAlY+EB-PVD TBC (vacuum annealed,after 1463 cycles at 1100 °C). White arrows indicate hafnium oxides.

454 U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

Pertinent characteristics of all MCrAlY coatings on Rene 142 arei) presence of hafnia particles in the TGO and ii) a very wavy interfacebetween TGO and bond coat, accompanied by large variations in TGOthickness and apparently entrapped metallic bond coat particles thatare in reality a consequence of the three-dimensional undulations ofthe TGO and a 2D-cut through three-dimensional oxide fingers. Anexample of a VPS-MCrAlY is shown in Fig. 6 that also disclosessignificant large spinel oxides within the TGO. This TGO microstruc-ture is quite similar to the TGO of the EB-MCrAlY-X version. Obviouslythe source of hafnium doesn't matter, and diffusion of this elementtowards the surface is fast enough in both cases. Since the diffusiondistance is larger for the Hf-containing substrate Rene 142 comparedto an Hf-containing bond coat, it develops the hafnium-rich oxidesslightly later and not during vacuum annealing. Accordingly, theseoxides are found mainly in the lower TGO region, accompanied bylarger undulations TGO/bond coat, while in the Hf-containing bondcoat they are more frequently observed in the upper TGO region.Consequently, an off-plane mixed zone is missing here. Diffusion of Hfin MCrAlY is much faster than for Y, and hence the tendency to surfaceenrichment is higher for Hf compared to Y as shown in [20] byexperimental evidencing a sputtered Hf interlayer. As detailed inprevious papers [2,3,21], the fact that TGOs on samples with thelongest lifetimes showed a very rough BC/TGO interfaces does notmatch models that have been developed for TBC spallation suggestingthat imperfections in that area might be initiation sites for large scalebuckling and subsequent spallation of the TBC. In the case of MCrAlYon Rene 142 and EB-MCrAlY-X on CMSX-4, hafnia containing oxidepegs and the large undulations may act as crack stoppers since aninterface crack has to change its propagation direction repeatedly if itfollows intimately the BC/TGO interface that possesses the lowestinterface toughness. The TGO is tightly bonded to the metallic bondcoat by these pegs, with a positive effect of the rare earth elementhafnium. In [22] plasma-sprayed NiCoCrAlY-SiHf+EB-PVD top coatsystems developed similar oxide pegs rich in Hf an Y suggesting thatpresence of mainly Hf always leads to oxide pegs and a typicalundulated TGO bond coat interface. Same observations have beendescribed for Hf-doped PtAl coatings [23]. Surprisingly, a similarundulated interface containing oxide pegs was observed for the VPS-MCrAlY systems on the Hf-free alloys despite of only traces of Hf in thebond coat. The TGO of version PtAl on Rene 142 shows only limitedsigns of rumpling on this alloy but it contains also large Hf-oxide pegsat the TGO bond coat interface [10].

3.4. Activation energy of cyclic lifetime data

Activation energies can provide indirect evidence to decode TBCfailure mechanisms. If compared with activation energies of basic

processes such as diffusion, oxide growth, phase transformation, crackpropagation etc., they can indicate potential mechanisms responsiblefor failure [20,25]. Furthermore, if the activation energy changes overa larger temperature interval of e.g. 1000–1100–1200 °C, this is anindication of a change in failure mechanism.

Although the current data are by far not robust due to the smallsample numbers, which immediately gets evident if the previousnovel bond coat data of 483 kJ mol−1 based on two failed samples [20]and the latest 424 kJ mol−1 based on meanwhile three tested samplesare compared, the following conclusions can be drawn.

The lowest activation energy achieved with PtAl on CMSX-4suggests that this system degrades slowly if the temperature is raised.The calculated activation energy is consistent with 225 kJ mol−1 fromdata given for PtAl on Rene N5 between 1100 and 1150 °C in [24],whereas other investigations for PtAl systems revealed higheractivation energies of 356 to 520 kJ mol−1 especially for thetemperature interval 1100 to 1200 °C [9,20]. The apparent failuremechanism of rumpling of the CMSX-4+PtAl system investigated hereseem to have a low activation energy. Intermediate values ofactivation energies in the range of 350 to 390 kJ mol−1 compare wellto inward-growing oxidation of alumina-forming alloys. The currentvalues for EB-PVD NiCoCrAlY and PtAl on IN100 fall into that regimeand agreewell with literature data that have been put together in [20].

The highest activation energy of 424 kJ mol−1 was found for thechemically modified EB-MCrAlY-X bond coat system. The arrange-ment of cracks may play a decisive role for failure, as derived from themicrographs in Fig. 5 and detailed in [20]. The high activation energyindicates a large temperature dependence of the failure mechanismand hence a large drop in cyclic life if the temperature is raised. But thedramatic gain in lifetimes still outrules the deficiencies envisaged athigh temperatures. Interestingly, a similarly high activation energy of418 kJ mol−1 can also be determined for an MCrAlY+EB-PVD TBCsystem given in [26] in the temperature range of 1150 to 1190 °C.Finally, the potential of a coating system can be realized more reliablyin view of associating absolute cyclic life and activation energy data.

4. Conclusions

Cyclic oxidation experiments of 7YSZ EB-PVD top coats on varioussubstrate–bond coat combinations revealed the followingconclusions.

– The longest lifetimes were achieved on a Hf-containing EB-PVDbond coat on CMSX-4 followed by coatings on the Hf-contain-ing substrate superalloy Rene 142. Whenever hafnium isinvolved, hafnia as well as oxide pegs are present in the TGOleading to large undulations at the interface TGO bond coat thatprolong TBC life. Local spinels in the TGO are not harmful forTBC life.

– The substrate alloy plays a decisive role for TBC life. There is nouniversal bond coat for all substrates, instead an optimized bondcoat treatment tailored for each system offers potential for life timeimprovements. Treatments such as annealing of the bond coat orsurface smoothing can prolong TBC life.

– Analysis of activation energies offers a valuable information toascertain possible failure mechanisms.

Acknowledgements

The authors gratefully acknowledge careful manufacture andtesting of the coatings by J. Brien, C. Kröder, H. Mangers, andD. Peters. O. Bernardi, W. Braue, H. Lau, U. Kaden, and B. Baufeldcontributed some of the TGO and lifetime investigations. Theprovision of the VPS bond coats and several discussions withR. Vaßen from FZ Juelich is gratefully acknowledged. Special thanksto M. Peters for reviewing the manuscript.

455U. Schulz et al. / Surface & Coatings Technology 203 (2008) 449–455

References

[1] U. Kaden, C. Leyens, M. Peters, W.A. Kaysser, in: J.M. Hampikian, N.B. Dahotre(Eds.), Elevated Temperature Coatings: Science and Technology III, 1999, p. 27.

[2] U. Schulz, M. Menzebach, C. Leyens, Y.Q. Yang, Surf. Coat. Technol. 146–147 (2001)117.

[3] U. Schulz, H. Lau, H.-J.R. -Scheibe, W.A. Kaysser, Zeitschrift f. Metallkd. 94 (2003)649.

[4] U. Schulz, K. Fritscher, in K. Kokini (ed.) ASME, New York, 1993, p. 163–172.[5] T.J. Nijdam, W.G. Sloof, Surf. Coat. Technol. 201 (2006) 3894.[6] U. Schulz, O. Bernardi, A. Ebach-Stahl, R. Vassen, Surface and Coatings Technology

(in press), doi:10.1016/j.surfcoat.2008.08.071.[7] M. Matsumoto, Surf. Coat. Technol. 202 (2008) 2743.[8] H.M. Tawancy, N. Sridhar, N.M. Abbas, D. Rickerby, J. Mater. Sci. 35 (2000) 3615.[9] N.M. Yanar, G. Kim, S. Hamano, F.S. Pettit, G.H. Meier, Mater high temp. 20 (2003)

495.[10] H. Lau, C. Leyens, U. Schulz, C. Friedrich, Surf. Coat. Technol. 165 (2003) 217.[11] G.M. Kim, N.M. Yanar, E.N. Hewitt, F.S. Pettit, G.H. Meier, Scri. Mater 46 (2002) 489.[12] U. Kaden, Phd thesis, Werkstoffwissenschaftliche Schriftenreihe, Vol. 56. 2003,

RWTH Aachen. ISBN 3-86130-094-X.[13] B. Baufeld, U. Schulz, Surf. Coat. Technol. 201 (2006) 2667.[14] O. Bernardie, diploma thesis, 2005, TU Darmstadt. p. 62.

[15] W. Braue, K. Fritscher, U. Schulz, C. Leyens, R. Wirth, Mat. Sci. Forum 461–464(2004) 899.

[16] W. Braue, U. Schulz, K. Fritscher, C. Leyens, R. Wirth, High Temp. Mater. 22 (2005)393.

[17] V.K. Tolpygo, D.R. Clarke, Surf. Coat. Technol. 200 (2005) 1276.[18] M.J. Stiger, Metall. Mater. Trans. A 38 (2007) 848.[19] C.G. Levi, E. Sommer, S.G. Terry, A. Catanoiu, M. Rühle, J. Am. Ceram. Soc. 86 (2003)

676.[20] K. Fritscher, U. Schulz, C. Leyens, Mater.wiss. Werkst.tech. 38 (2007) 734.[21] U. Schulz, K. Fritscher, and W.A.Kaysser, COST 2002, Liege, J. Lecomte-Beckers, M.

Carton, F. Schubert, P.J. Ennis, 2002, p. 483–492.[22] C. Mercer, S. Faulhaber, N. Yao, K. McIlwrath, O. Fabrichnaya, Surf. Coat. Technol.

201 (2006) 1495.[23] J. Nesbitt, B. Nagaraj, J. Williams, in: N.B. Dahotre, J.M. Hampikian, J.E. Morral (Eds.),

Elevated Temperature Coatings: Science and Technology IV, TMS, 2001, p. 77.[24] M.J. Stiger, N.M. Yanar, M.G. Topping, F.S. Pettit, G.H. Meier, Z. Met.k.d. 90 (1999)

1069.[25] R.M. German, Sintering Theory and Practice, Wiley, 1996.[26] D.S. Rickerby, Advanced Coatings for High Temperatures, Nice, Forum of

Technology, 2002.