tem characterization of plasma-sprayed thermal barrier coatings and ceramic–metal interfaces after...

Ž .Thin Solid Films 301 1997 105–114

TEM characterization of plasma-sprayed thermal barrier coatings andceramic–metal interfaces after hot isostatic pressing

H.C. Chen, J. Heberlein, E. Pfender )

Department of Mechanical Engineering, UniÕersity of Minnesota, 111 Church Street S.E., Minneapolis, MN 55455, USA

Received 16 September 1996; accepted 12 December 1996

Abstract

Ž . Ž .Plasma-sprayed thermal barrier coatings TBCs consisting of both a Ni–5%wt. Al bond coat Ni–Al and a thermally insulatingŽ .8%wt. Y O partially stabilized ZrO ceramic top coat were subjected to hot isostatic pressing HIP . Transmission electron microscopy2 3 2

Ž .TEM with qualitative microanalysis was utilized to characterize the microstructural changes occurring during HIP treatment and tow Ž .xclarify the nature of the ceramic–metal interface ZrO – Ni–Al both in the as-sprayed and hipped states. During HIP treatment, the2

as-sprayed ZrO experienced a certain degree of hot plastic deformation and recrystallization, forming dense and crackless fine-equiaxed2

tetragonal grains. However, due to nonuniformities of the as-sprayed structure, the degree of recrystallization of the original columnargrains changed from location to location, resulting in a recrystallized microstructure ranging from coarse equiaxed grains 1–2 mm in sizeto co-existing grains of equiaxed and columnar shapes, including still intact columnar grains. Through TEM observations of a

Ž .plasma-sprayed ZrO – Ni–Al interface more than 20 mm long, it was found that the interface is amorphous in nature and unevenly thick2

and even discontinuous in some locations. The amorphous interfacial layer is composed largely of the Ni–Al bond coat materials togetherŽ .with some ZrO ceramic material. After HIP treatment, the amorphous ZrO – Ni–Al interfacial layer is replaced by a continuous, dense2 2

and polycrystalline Al O layer which can act as a diffusion barrier for elements both from the ZrO coating to the bond coat and to the2 3 2

substrate, and also in the opposite direction.

Ž .Keywords: Ceramics; Coatings; Plasma processing and deposition; Transmission electron microscopy TEM

1. Introduction

Ž .Plasma-sprayed thermal barrier coatings TBCs areincreasingly used for thermal insulation of hot stage gasturbine components to improve engine efficiency and ser-vice life. Typical state-of-the-art TBCs use 8%wt. Y O2 3

partially stabilized ZrO , which produces a nonequilibrium2Ž X .tetragonal phase t -phase with high-yttria content during

w xplasma spraying due to rapid cooling 1 of the spraymaterial. Recently, studies have been reported on dealingwith the effects of chemical and phase composition on

w xproperties of the TBCs 2,3 , on the microstructure both inw x w xthe as-sprayed 4,5 and annealed states 6 , and on failure

w xmechanisms 8,9 of the TBCs and post-treatment appliedw xto the TBCs 10 .

The TBCs may easily fail because of the poor bondingbetween the ceramic coating and the bond coat and thehigh porosity existing in the ceramic coating. It has beenshown that cracks are preferentially initiated and propagate

) Corresponding author.

along the ceramic–metal interface, which later caused thew xceramic coating to spall 9 . Therefore, the nature and

bonding mechanism of the ceramic–metal interface is akey factor in determining the service life of the TBCs.Unfortunately, little attention has so far been paid to themicrostructure of the ceramic–metal interface, especiallyto characterization by TEM. There are only a few refer-ences available concerning microstructures of TBCs in theas-sprayed or annealed states and of ceramic–metal inter-

w xfaces under oxidizing conditions 11 . In this paper, TEMtechniques were utilized to clarify the nature of the ce-ramic–metal interface, both in the as-sprayed and hippedstates, and to characterize microstructural changes of theceramic coating during HIP treatment. Other changes, suchas effects on porosity and fracture during HIP treatment,

w xhave been described elsewhere 10 .

2. Experimental method

The commercial 8%wt. Y O partially stabilized ZrO2 3 2Žpowder AI 1075, Alloy International, Miller Thermal,

0040-6090r97r$17.00 Copyright q 1997 Elsevier Science B.V. All rights reserved.Ž .PII S0040-6090 96 09604-6

( )H.C. Chen et al.rThin Solid Films 301 1997 105–114106

Fig. 1. A series of micrographs showing the nature of lamellae in the as–sprayed ZrO coating. Note that columnar grains are well developed in the2Ž . Ž . Ž .lamellae and pores are visible between the lamellae of the coating shown by arrows . a SEM morphology of the cross-sectional coating. b SEM

Ž .morphology of fracture surface of the coating. c TEM morphology of the cross-sectional coating.

.Inc., USA used in these experiments was spherical inshape with a size of y200 to q320 mesh. The substratewas a 15 mm diameter and 40 mm long cylinder consistingof cold rolled mild steel containing 0.18%wt. carbon. Thesubstrate was ultrasonically cleaned and grit-blasted with60 mesh Al O under pressures of around 0.4 MPa for2 3

Žapproximately 30 s. A 0.1 mm thick Ni–Al s95–5 %wt.,.AI 1037, Alloy International, Miller Thermal, Inc., USA

bond coat followed by a 0.3 mm thick ZrO was then2

plasma-sprayed on the surfaces of the rotating samplesusing a Miller SG-100 plasma spray torch. During spray-ing, the substrate was cooled with forced air. The Miller

Ž . Ž .Fig. 2. TEM micrographs showing small amounts of lenticular martensite existing in the as-sprayed ZrO coating shown by arrows . a Bright field2Ž .image of the martensitic transformation zone. b Dark field image of the martensitic transformation zone.

( )H.C. Chen et al.rThin Solid Films 301 1997 105–114 107

Fig. 3. Typical SEM micrograph of the cross-sectional ZrO coating after2

hipping treatment, showing a dense and crackless coating.

SG-100 plasma spray torch used in these experiments wasŽoperated in the subsonic mode anode 2083-175, cathode

.1083A-129, gas injector 2083-113 . The spray powder,suspended in an argon carrier gas, was injected through the

Žinternal powder feed port located about 10 mm upstream

Table 1Plasma spraying conditions

Parameters PowderNirAl ZrO2

Ž .Arc current A 800 850Ž .Primary gas flow rate Ar, slm 45 41.5Ž .Secondary gas flow rate He, slm 16 16

Ž .Carrier gas flow rate Ar, slm 4.5 4.5Ž .powder feed rate grmin 39 25Ž .Standoff distance mm 120 90

Ž .Substrate rotation speed rpm 60 60Ž .Torch nozzle i.d. mm 8 8

Ž .Powder injector tube i.d. mm 2 2

. y1of the torch exit with a velocity of about 7 m s . Theplasma spraying conditions for this experiment are summa-rized in Table 1.

Ž .Hot isostatic pressing HIP treatment of the coatingwas performed in a Paterson Instrument. The sprayedsample was raised to a temperature of 1250 8C over 2 h inan argon environment, and then an isostatic pressure of200 MPa was applied at this temperature for 8 h beforefurnace-cooling.

Ž .Fig. 4. Typical TEM micrographs of the hipped ZrO coating showing that the microstructure consists of fine equiaxed grains. a Morphologies of the fine2Ž . Ž .equiaxed grains implies that the as-sprayed coating underwent complete recrystallization after HIP treatment. b Higher magnification of a showing

Ž . Ž .twins existing within some of the equiaxed grains shown by arrow . c SAED pattern of the equiaxed grains showing the tetragonal structure of ZrO .2


Ž .The post-treated and untreated as-sprayed specimenfor metallographic examination were first impregnated withepoxy resin in a vacuum jar and then cut perpendicular tothe coating surface with a low-speed saw, and then pol-ished with a Struer’s grind–polish system. Microstructureand fracture surface morphologies of the cross-section ofthe coating after controlled fracture of the coating–sub-

w x Ž .strate system 10 were analyzed by SEM JEOL 840-P .Techniques for preparation of TEM cross-sectional

samples of the coating–substrate interface are similar tow xthose described in Ref. 12 . The cylindrical sample was

first mounted with epoxy in a stainless steel tube withinner diameter of 16 mm, and then a thin slice 1 mm thickwas sectioned along the direction normal to the coatingsurface. The slice was then attached to a tripod polisherwith epoxy and thinned mechanically and polished fromboth sides until the ceramic coating became transparent.Finally the sample thickness has been further reduced with

Ža low-angle ion-thinning machine Precise Ion Polish Sys-.tem . The samples containing the coating–substrate inter-

face were examined and analyzed using the Philips EM-30transmission electron microscope, equipped with a qualita-tive microanalysis unit.

3. Experimental results and analysis

3.1. Characterization of the hipped ZrO coating2

Plasma-sprayed coatings have been described as abuild-up of flattened splats, having lamellar structures. As

w Ž .xshown in a cross-section of the coating Fig. 1 a , theplasma-sprayed coating shows some cracks and pores. Thebuilding block nature of the coating structure and thecolumnar grain features within a splat can be seen on thefracture surface of a coating after the controlled fracturew x w Ž .x10 of the coating-substrate system Fig. 1 b . By usingTEM, pores between the lamellae or between splats wereidentified to be caused by poor contact between them, as

Ž .shown in Fig. 1 c . Previous work showed that plasma-sprayed ZrO coatings consist entirely of high-yttria con-2

w xtent and nontransformable tetragonal phases 10 , but TEMobservations in this work revealed some small martensitictransformation zones existing in the coating, as shown byarrows in Fig. 2. The martensite appeared in the lenticularmorphology. This discrepancy might be attributed to thefact that the grinding and ion-milling employed in the foil

Ž .preparation might have induced stress or strain high

Ž .Fig. 5. TEM micrographs showing different degrees of recrystallization occurring in the different area of the ZrO coating during HIP treatment. a Zone2Ž .of complete recrystallization and grain growth with the grain size of 1–2 mm. b Partial recrystallization zone, with the fine equiaxed grains and columnar

Ž .grains coexisting in lamellae 1, 2 and 3. c Mixed zone of complete and partial recrystallizations. Note coarse equiaxed grains in lamella 5 and partialŽ .recrystallization in lamellae 4 and 6. d Nonrecrystallization zone where columnar grains remained.


enough to produce the martensitic transformation, or alter-natively that the X-ray diffractometers employed in theprevious work might not be sensitive enough to detect thesmall amount of martensite that was formed in the

w xplasma-spraying process 6 .Ž .As indicated in a cross-section of the coating Fig. 3 ,

after HIP treatment the ZrO coating experienced substan-2

tial changes, from the porous and micro-cracked morphol-ogy in the as-sprayed state to the dense and cracklessmorphology with micro-voids less than 0.1 mm in size.Fine microstructural features of the hipped ZrO coating2

were characterized by TEM. Due to hot plastic deforma-tion and recrystallization, the hipped coating consisted of aseries of fine equiaxed grains with sizes ranging from 0.1

Ž .to 0.6 mm, as shown in Fig. 4 a . Under higher magnifica-tion, it was found that twins exist in some equiaxed grainsw Ž .x Ž .Fig. 4 b . A selected area electron diffraction SAEDpattern showed the equiaxed grain exhibited the tetragonal

w Ž .xstructure Fig. 4 c . However, due to nonuniformities ofthe as-sprayed coating, the driving force for recrystalliza-

Ž .tion diffusion for atoms changed from location to loca-tion, representing several zones with different degree ofrecrystallization existing in the coating, as shown in Fig. 5.In some regions, the as-sprayed columnar structure under-went not only complete recrystallization, but also graingrowth, forming coarse equiaxed grains 1–2 mm in sizew Ž .xFig. 5 a ; while in other regions the original columnargrains remained intact, corresponding to the nonrecrystal-

w Ž .xlization zone Fig. 5 d . There are some transition zones,Ž .exhibiting either a partial incomplete recrystallization

w Ž .xzone Fig. 5 b or a mixed zone of complete and incom-w Ž .xplete recrystallization Fig. 5 c . Each of these zones

contains either the co-existing grains of equiaxed andw Ž .xcolumnar shapes lamellae 1, 2 and 3 of Fig. 5 b or the

w Ž .xmixed form of the coarse grains lamella 5 of Fig. 5 cwand the above-mentioned co-existing grains lamellae 4

Ž .xand 6 of Fig. 5 c .In addition, in the recrystallization zone, TEM also

Ž .indicated some stacking faults Fig. 6 , which are thoughtto be caused by isostatic pressing applied to the coating

Fig. 6. TEM micrograph showing stack faults existing in the recrystalliza-tion zone of the hipped ZrO coating.2

Ž .Fig. 7. TEM cross-section of the as-sprayed ZrO – Ni–Al interface. The2

SAED pattern, a series of faint rings, shows the amorphous nature of theinterfacial region. The amorphous interfacial layer is unevenly thick from0.1–0.8 mm and discontinuous in some locations. Also visible are some

Ž .pores in the interface shown by arrows .

leading to the deformation of the newly formed equiaxedgrains.

3.2. Plasma-sprayed ceramic–metal interface

Although the general technique for foil preparationallowed thinning of the ceramic coating and of the metal


Ž .Fig. 8. a Close-up of the amorphous interfacial region indicated by the arrow A in Fig. 7. Note that the rough amorphous morphology may be caused byŽ .ion-milling. b EDS of the ZrO coating adjacent to the amorphous interfacial layer, showing that a small amount of Ni diffusion occurred from the Ni–Al2

Ž . Ž .bond coat into the ZrO coating. c EDS of the amorphous interfacial layer, showing high Ni content with trace amounts of Zr, Y. d EDS of the Ni–Al2

bond coat adjacent to the amorphous interfacial layer, showing the presence of Al, which is related to the composition of the bond coat.

Ž .inner layer bond coat as well as of the interfacial regionbetween them, the interfacial region cannot always beimaged under the TEM, because it may have been prefer-entially etched away during ion-milling when the ceramiccoating and the bond coat were thin enough for imaging.Therefore, in order to examine and analyze the microstruc-

ture and composition of the interfacial region, the ion-mill-ing process has to be stopped before a visible hole devel-ops, which usually provides some thinned regimes forexamining the structures of the ceramic coating and bondcoat. Fig. 7 shows such an ‘‘incomplete ion-milling’’ foil

wwhere more than a 20 mm long ceramic–metal ZrO –2

Ž . Ž . Ž .Fig. 9. a TEM micrograph shows that the as-sprayed ZrO – Ni–Al interface is bordered on the ZrO side by a thin layer of ultra-fine grains. b SAED2 2

pattern shows that the ultra-fine grains exhibit a cubic structure.


Ž .xNi–Al interfacial line was imaged. However, the ce-Ž . Ž .ramic coating ZrO and bond coat Ni–Al at both sides2

Ž .of the interface remain dark unable to image under TEMbright field illumination. A series of TEM micrographs inFig. 7 show that the plasma-sprayed ZrO was not me-2

chanically separated from Ni–Al. There exists an interfa-cial layer between them. The SAED pattern, a series offaint rings, shows the amorphous nature of this interfacial

Ž .layer Fig. 7 . The amorphous interfacial layer was foundto be unevenly thick: from less than 0.1–0.8 mm and evendiscontinuous in some locations. Some stretched pores inthe interface may be due to preferential erosion duringion-milling. Qualitative microanalysis conducted simulta-neously with TEM observations revealed changes in thechemical composition across the interfacial layer, as shown

Ž .in Fig. 8. Fig. 8 a shows a close-up of the amorphousinterfacial region indicated by the arrow A in Fig. 7, wherethe rough morphology in the amorphous interfacial phasemay be caused by ion-milling. EDS of the ZrO coating2

Ž .adjacent to the interfacial layer is shown in Fig. 8 b ,showing small amounts of Ni diffusion from the bond coatto the ceramic coating. It is interesting to notice that theinterfacial layer is rich in Ni with trace amounts of Zr and

w Ž .xY Fig. 8 c , which implies that the interfacial layercontains large amounts of the bond coat material togetherwith some ceramic material. EDS of the bond coat adja-cent to the interfacial layer, on the other hand, reveals noZr and Y but Al, which relates to the chemical composi-tion of the bond coat.

Moreover, TEM also showed that the plasma-sprayedŽ .ZrO – Ni–Al interface was bordered on the ZrO side by2 2

a thin layer of ultra-fine grains, which have been identifiedby SAED to be of a cubic structure, as shown in Fig. 9.

3.3. Hipped ceramic–metal interface

Ž .The plasma-sprayed ZrO – Ni–Al interface changed2

substantially when the coating was subjected to HIP treat-ment. Fig. 10 shows a TEM low-magnification micrograph

Ž .of the hipped ZrO – Ni–Al interface. It seems that the2

interfacial layer becomes thicker and continuous after HIPtreatment. The composition and structure of this hippedinterfacial layer is shown in Fig. 11. A close-up in Fig.Ž .11 a shows that the amorphous interfacial layer in the

as-sprayed state has been replaced by a dense and poly-crystalline layer due to diffusion of atoms leading tocrystallization and chemical reaction at the interface under

wthe condition of HIP treatment. The SEAD pattern Fig.Ž .x w Ž .x11 b and EDS Fig. 11 d show that the hiped interfacial

Ž .layer is composed of hexagonal close-packed hcp a-Al O . From EDS at different locations across the Al O2 3 2 3

layer, it follows that the dense Al O layer can act as a2 3

diffusion barrier for Fe from the substrate through thebond coat to the ZrO coating, because the Fe contents of2

w Ž .xthe Al O layer Fig. 11 d and ZrO coating adjacent to2 3 2w Ž .xthe interfacial layer Fig. 11 c are much lower than that

Ž .Fig. 10. Typical TEM cross-section of the hipped ZrO – Ni–Al inter-2

face. Note that a continuous but unevenly thick interfacial layer wasformed between the ZrO coating and the Ni–Al bond coat.2

wof the bond coat adjacent to the interfacial layer Fig.Ž .x11 e .The Al O phase in the hipped interface is bordered2 3

with ZrO and Ni–Al of different morphologies, as shown2

in Fig. 12. Although the interface is dense across itsthickness, the grain size changes from the Al O –ZrO2 3 2

Ž .interface to the Al O – Ni–Al interface. At the Al O –2 3 2 3

ZrO interface, Al O has a microstructure consisting of2 2 3w Ž .xultra-fine equiaxed grains about 90 nm in size Fig. 12 c ,

Ž .while at the Al O – Ni–Al interface, the grain size of2 3w Ž .xAl O increases to about 0.4 mm Fig. 12 a . Note that the2 3Ž .arrowed laminar pore in Fig. 12 a again is accounted for

by preferential erosion during ion-milling. Under highŽ .magnification, TEM reveals an ultra-thin about 25 nm

and high-Fe content layer at some locations of the Al O –2 3Ž .Ni–Al interface. This ultra-thin layer cannot be identifiedby TEM due to its small volume. It may consist of a solidsolution or compound containing Ni and Fe. In addition, atthe Al O –ZrO interface, dislocations can be found which2 3 2

are believed to be caused by the release of thermal stressarising from the Al O growth.2 3

4. Discussion

4.1. Recrystallization of the plasma-sprayed coating dur-ing HIP treatment

It has been shown that the plasma-sprayed coatingexperienced a certain degree of hot plastic deformation andthe columnar grains recrystallized to form equiaxed grainswhen the coating is subjected to HIP treatment. Due tononuniformity of the plasma-sprayed structure, resultingfrom nonuniform cooling during plasma-spraying, the grainsize and morphology of the recrystallized microstructurechanges from location to location. This is attributed to thechange in the driving force for recrystallization, leading todifferent degrees of recrystallization in different areas ofthe coating.


During HIP treatment, there are two possibilities forcausing recrystallization of the plasma-sprayed structurew x10 , i.e. static recrystallization and dynamic recrystalliza-tion. The static recrystallization takes place because thecolumnar grains of the plasma-sprayed coating aremetastable and in a high energy state, which tends toreturn to the thermodynamic equilibrium state by formingfine-equiaxed grains which minimizes the grain boundaryenergy of the system. The dynamic recrystallization, how-ever, takes place due to hot plastic deformation, whichcauses the columnar grains to be broken up into pieces,

forming many sub-structures and crystal defects withingrains. During plasma-spraying, the cooling conditions ofthe molten splats change from location to location, depend-

w xing on a number of factors 13 , such as droplet tempera-ture, viscosity, impacting velocity and contact between thesplat and the previously deposited material. In areas wherethe cooling rate is larger, finer and shorter columnar grainswithin the splat are formed, while in areas where thecooling rate is smaller, coarser and longer columnar grainsare formed. Fine and short grains contain high grainboundary energy which provides high driving forces for

Ž .Fig. 11. Structure and compositions of the interfacial Al O layer in the hipped ZrO – Ni–Al interface. Note that the Al O layer can act as a diffusion2 3 2 2 3Ž .barrier for Fe from the substrate through the bond coat into the ZrO coating. a TEM micrograph of the interfacial Al O layer shows the dense and2 2 3

Ž . Ž . Ž . Ž .polycrystalline nature. b SAED pattern of the interfacial layer showing a-Al O structure. c , d and e EDS of the ZrO coating adjacent to the2 3 2

interfacial layer, of interfacial layer, and of the Ni–Al bond coat adjacent to the interfacial layer. Note that there is a limited amount of Fe existing in theinterfacial layer and the ZrO coating.2


recrystallization. Moreover, they are easy to be broken intopieces during HIP treatment. Therefore, in these areas thestatic and dynamic recrystallization processes proceed tocompletion, forming equiaxed grains. In contrast, the coarseand long columnar grains make the recrystallization pro-cesses difficult and incomplete, and thus the co-existinggrains of the equiaxed and columnar shapes are formed,including intact columnar grains. It is possible that thegrain growth takes place in the newly formed equiaxedgrains or in the local original equiaxed grains which wereformed at the periphery of the splat or within highlyundercooled splats during plasma-spraying.

4.2. The nature of the plasma-sprayed ceramic–metalinterface

According to the TEM observations and EDS analysis,Ž .the plasma-sprayed ZrO – Ni–Al interface was composed2

of the amorphous phase containing large amounts of thebond coat material together with some ceramic coatingmaterial. This is believed to be a consequence of theplasma-spraying process. When plasma-spraying ZrO ce-2

ramics, high temperatures in the plasma are needed in

order to melt the particles. In particular, the first particlesarriving at the bond coat have to be heated to temperatures

Ž .exceeding the melting point 2750 8C to form ‘‘super-heated’’ droplets which stick on the surface of the bondcoat when impacting. On impact, heat transfer from the‘‘superheated’’ ZrO droplets to the relatively low melting2

Ž .point 1500 8C of the Ni–Al bond coat will cause partialmelting of the outer surface of the Ni–Al layer, forming athin layer of ‘‘mixed liquid’’ composed largely of theNi–Al bond coat material together with some ZrO mate-2

rial. In comparison with the entire coating, this thin layerof ‘‘mixed liquid’’ is very small in volume. When it

Ž .spreads over the cool substrate bond coat , rapid quench-ing occurs, resulting in the formation of an amorphousphase. Because the ‘‘mixed liquid’’ fuses the ZrO coating2

and the bond coat together upon cooling, it is expected thatthis amorphous interfacial layer has a positive effect onadhesion of the ceramic coating to the bond coat. This

w xfinding is similar to that of Harmsworth et al. 5 .In the region adjacent to the amorphous interfacial

layer, the ZrO microstructure consists of fine cubic2

equiaxed grains due to the reduced cooling rate. Thisimplies that the transformation of the cubic structure totetragonal structure was suppressed by the cooling.

Ž . Ž .Fig. 12. TEM micrographs show some features of Al O layer in the hipped ZrO – Ni–Al interface. a Dense equiaxed grains about 0.4 mm in size were2 3 2Ž .generally found in the area close the Al O – Ni–Al interface. Note that the arrowed stretched pore is accounted for by preferential erosion during2 3

Ž . Ž . Ž .ion-milling. b High magnification shows an ultra-thin layer about 25 nm thick with high Fe content existing in some locations of the Al O – Ni–Al2 3Ž . Ž . Ž .interface shown by arrow . c Dense and ultra-fine equiaxed grains about 90 nm in size existing in the Al O –ZrO interface. d High magnification of2 3 2

Ž .the region arrowed in c shows dislocations existing in the Al O side of the Al O –ZrO interface.2 3 2 3 2


4.3. The nature of the hipped ceramic–metal interface

Due to the fact that HIP treatment creates an environ-ment which is conducive to atomic diffusion, substantialchanges take place not only within the ZrO coating, but2

Ž .also in the ZrO – Ni–Al interface during HIP treatment.2Ž .After HIP treatment, the amorphous ZrO – Ni–Al inter-2

facial layer is replaced by a continuous and dense andpolycrystalline Al O layer. The grain size of the Al O2 3 2 3

phase at the hipped interface is larger at the Ni–Al sidethan at the ZrO side. This implies that the Al O phase2 2 3

nucleates at the Al O –ZrO interface and then grows. As2 3 2

the Al O layer thickens, additional thermal expansion2 3

occurs between the ZrO coating and the Al O phase,2 2 3

leading to thermal stresses which are responsible for dislo-cations observed at the Al O side of the Al O –ZrO2 3 2 3 2

interface. Nucleation and growth of Al O at the Al O –2 3 2 3

ZrO interface requires diffusion of Al from the bond coat2

across the formed Al O layer to the ZrO coating. Al-2 3 2

though growth of Al O consumes some parts of Al in the2 3Žbond coat, it prevents oxygen coming from the oxidizing

.atmosphere entering the bond coat. According to theresults of EDS, the dense Al O layer can also act as a2 3

diffusion barrier for Fe diffusing from the substrate acrossthe bond coat to the ZrO coating. It should be pointed out2

that the Al O interfacial layers so formed under HIP2 3

treatment are dense across their thickness, which is differ-w xent from the results of Leleit et al. 7 who reported that

the Al O interfacial layer formed at the annealed ZrO –2 3 2

NiCrAlY interface which was dense at the Al O –ZrO2 3 2

interface, but porous at the Al O –NiCrAlY interface.2 3

This difference may be due to the compaction effect ofHIP treatment which closed up the pores of the Al O2 3

interfacial layer.

5. Conclusions

During HIP treatment, the as-sprayed ZrO coating2

experienced a certain degree of hot plastic deformation andrecrystallization, forming a dense and crackless layer offine-equiaxed tetragonal grains. However, due to nonuni-formity of the as-sprayed structure, the degree of recrystal-lization of the original columnar grains changed fromlocation to location, resulting in a recrystallized mi-crostructure ranging from coarse equiaxed grains 1–2 mm

Ž .in size complete recrystallization and grain growth zoneto co-existing grains of equiaxed and columnar shapesŽ .partial or incomplete recrystallization zone and to intact

Ž .columnar grains nonrecrystallization zone .Ž .Plasma-sprayed ZrO – Ni–Al interfaces were found to2

consist of a layer of an amorphous phase which wascomposed largely of the Ni–Al bond coat materials to-gether with some ZrO ceramic material. The amorphous2

interfacial layer was unevenly thick and even discontinu-ous in some locations. After HIP treatment, the amorphous

Ž .ZrO – Ni–Al interfacial layer was replaced by a continu-2

ous, dense and polycrystalline Al O layer which can act2 3

as a diffusion barrier for elements both from the ZrO2

coating to the bond coat and to the substrate and also fromthe opposite direction. It is expected that this dense Al O2 3

layer has a positive effect on the service life of the TBCs.

Acknowledgements

This work has been supported by NSF through the ERCfor Plasma-Aided Manufacturing, grant EEC-87-21545.The government has certain rights in this material.

References

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