Thermal Barrier Coatings for the 21st Century

Transcription

1 Thermal Barrier Coatings for the 21st Century M. J. Stiger, N. M. Yanar, M. G. Topping, F. S. Pettit, and G. H. Meier Department of Materials Science and Engineering University of Pittsburgh Pittsburgh, PA ABSTRACT Yttria stabilized zirconia (YSZ) thermal barrier coatings (TBCs) fabricated via electron beam physical vapor deposition (EBPVD) provide some unique properties for aerofoil applications. Such coatings are usually deposited on diffusion aluminide or MCrAlY bond coats on superalloy substrates. During deposition of the YSZ-TBC on the bond coat, a thermally grown oxide (TGO) consisting primarily of -Al 2 O 3 is formed between the TBC and bond coat. The lives of these TBCs in oxidizing environments is determined by the interplay of the stored elastic energy driving spallation versus the interfacial toughness of the TGO/bond coat interface since failure of the TBC system often occurs at this interface. The microstructures of EBPVD-TBCs in the as-processed and exposed conditions have been documented using a variety of techniques to determine and describe the failure mechanisms. It is shown that the failure mechanisms for TBC s are dependent on the bond coat fabrication procedure and the exposure environment. Finally, possible schemes for improving TBC systems are outlined. INTRODUCTION Thermal Barrier Coatings The earliest thermal barrier coatings (TBCs) were frit enamels which were applied to aircraft engine components in the 1950s [1]. The first ceramic TBCs were applied by flame spraying and, subsequently, by plasma spraying. The ceramic materials were alumina and zirconia (MgO or CaO-stabilized), generally applied directly to the component surface. The effectiveness of these coatings was limited by the relatively high thermal conductivity of alumina and problems with destabilization of the zirconia-based materials [1]. Important developments included the introduction of NiCrAlY bond coats and plasma-sprayed Y 2 O 3 -stabilized zirconia topcoats in the mid-1970s and the development of electron beam physical vapor deposition to deposit the topcoat in the early 1980s [1]. These types of coatings have been used for many years on combustion liners but, with advanced thermal barrier coatings, vanes, and even the leading edges of blades, can now be coated. The use of thermal barrier coatings (TBCs) has resulted in a significant improvement in the efficiency of aircraft gas turbines [1-3] and has the potential to do the same in land-based gas turbines. The use of TBCs can achieve temperature differentials across the coating of as much as 175C. Figure 1 (After DeMasi- Marcin and Gupta [4]) shows the improved temperature capability of nickel-base superalloys over the years and the effective improvement associated with the use of TBCs 1

2 to lower the metal temperature. The current issue is the extent to which the properties of TBC systems can be improved further. Typical systems, shown schematically in Figure 2, consist of a nickel-base superalloy substrate coated with MCrAlY (M=Ni,Co) or a diffusion aluminide bond coat which forms an alumina layer (thermally-grown oxide, TGO) on to which is deposited a yttria-stabilized zirconia (YSZ) TBC. The TBC can be deposited by air plasma spraying (APS), or electron beam physical vapor deposition (EBPVD). The EBPVD coatings are used for the most demanding applications, such as the leading edges of airfoils. Figure 3 presents cross-section micrographs of typical APS and EB-PVD TBCs. The major challenge for the development of improved TBCs is coating durability, particularly the resistance of the coating to spalling. There are a number of degradation modes which can limit the life of a TBC and these must be understood in order to make lifetime predictions for existing systems and to provide the basis for the development of improved TBC systems. These include: Cracking within the ceramic layer which leads to spalling of part of the TBC. Cracking along the interface between the TGO and the bond coat which results in spalling of the entire TBC. Sintering of the TBC at the outer surface, where the temperature is highest, which increases the thermal conductivity of the TBC and can increase the total amount of elastic energy stored in the coating, which provides additional driving force for cracking and spalling of the coating. Particle erosion, which causes a continued wearing-away of the coating and, for large particles, can produce cracks in the coatings and along the interface between the TBC and the bond coat. Several factors are important with regard to cracking within the TBC or along the TBC-bond coat interface These include: the stress state in the zirconia layer, the microstructure of the bond coat, the thickness of the TGO, the stress state in the TGO, and the fracture resistance of the interface between the bond coat and the TGO. It was apparently first pointed out by Miller [5] and is now generally accepted that oxidation of the bond coat is a critical factor controlling the lives of EBPVD TBCs. It is now well established that the ability of a bond coat to form an -alumina layer with negligible transient oxidation and the adherence of the alumina to the bond coat are critical factors in controlling the durability of TBCs. Experience with aircraft engines has shown that bond coat oxidation and the ability to resist spalling of the TBC from the bond coat are critical factors determining coating life. The bond coats develop a thermally grown oxide (TGO) layer during fabrication. The TGO grows thicker during exposure of the TBC. It is generally observed for plasma-sprayed TBCs that bond coats which are "good alumina formers" provide the longest lives and result in fracture within the zirconia some distance above the bond coat [3, 6,7]. Bond coats which tend to produce significant amounts of spinel in the 2

3 oxide layer result in spallation at the zirconia/bond coat oxide interface and generally exhibit shorter lives. The EBPVD coatings are generally thought to spall at the TGO/bond coat interface or in the alumina layer [3]. Therefore, for the most advanced TBC systems, a critical problem is the growth of the TGO and its adherence to the bond coat The residual stress in the alumina layer on the bond coat also plays a significant role in spallation behavior. This residual stress is determined by the growth stresses and thermal stresses and any stress relaxation which has occurred by plastic deformation of the alloy and/or oxide. It is now also clear that residual stresses in the TBC topcoat can also play a role in spalling. The focus of this paper is the mechanisms of degradation of current state-of-the-art TBCs and how the technology may develop in the 21st Century for producing more durable coatings.. EXPERIMENTAL PROCEDURE The system which will be emphasized in this paper consists of the single crystal superalloy N5 coated with a platinum aluminide bond coat and an EBPVD 7wt% yttriastabilized zirconia TBC. The specimens are circular disks 25.4 mm in diameter and 3.2 mm thick. The specimens are coated on all surfaces with a nominally 50m thick bond coat and on one side only with a 100m thick TBC. Where appropriate for comparison purposes, results for an APS TBC with a NiCoCrAlY bond coat will also be presented. The specimens have been exposed to isothermal and cyclic oxidation at temperatures between 1100 and 1200C. Each cycle consisted of 45 minutes in the hot zone and 15 minutes in the cold zone. The times to spalling have been determined for the various exposures. The as-received and exposed specimens have been characterized by optical microscopy (OM), scanning electron microscopy (SEM), and cross-section transmission electron microscopy (XTEM). TBC Failure RESULTS AND DISCUSSION Figure 4 presents a macroscopic photograph of an EB-PVD TBC on a Ptaluminide bond coat which has failed after 1287 cycles at 1100C in dry air. The failure initiated at a specimen edge and propagated as an elongated buckle which branched several times. Figure 4 also shows an area of the bond coat from which the coating has spalled. Here the fracture path has been mainly along the TGO/bond coat interface, i.e. the fracture surface consists mainly of exposed bond coat. However, there are areas where the fracture path has entered the TGO and the YSZ. In fact, in some areas of the specimen, the fracture path traveled for significant distances along the TGO/YSZ interface. Figure 5 presents such an area where the underside of the spalled TBC has large areas where the YSZ is exposed. 3

4 Figure 6 indicates the effect of temperature on the time to failure. Here the inverse of the time at elevated temperature (0.75 times the number of cycles) is plotted versus reciprocal temperature. The data follow a reasonable straight line with failure times decreasing from about 1000 h at 1100C to 50 h at 1200C. Figure 6 also contains one failure time for a specimen exposed at 1100C in air with 0.1 atm water vapor. The presence of the water vapor decreases the failure time by roughly a factor of two. This result will be discussed subsequently. Factors Affecting TBC Failure TBC Sintering Sintering of the TBC has also been observed during high temperature exposure. Figure 7 compares the microstructure of an EB-PVD TBC in the as processed condition with that of an identical coating which has been exposed for 10 hours at 1200C. Sintering between the columns is evident. These observations are consistent with the dilatometry measurements of Zhu and Miller [8] who observed 0.1% shrinkage of cylindrical specimens of plasma-sprayed ZrO 2-8wt%Y 2 O 3 exposed for 15 hours at 1200C. Figure 8 illustrates the sintering behavior of an APS TBC. The as-processsed coating, Figure 8a., contains a substantial number of microcracks which have sintered closed after 100 h exposure at 1200C, Figure 8b. Experiments using a laser-imposed temperature gradient on a TBC coating have also resulted in substantial sintering in the outer regions of the coating [9]. Sintering has been observed to increase the effective Young s modulus of the TBC to values approaching that for dense zirconia (200GPa) [8, 10] and to result in residual stresses in the MPa range [10, 11]. These changes greatly increase the amount of elastic energy stored in the TBC and can enhance the driving force for scale spallation. Bond Coat Oxidation Figure 3 is an SEM cross-section showing the various layers in the coated specimens. The EB-PVD TBC is composed of large columnar grains of YSZ that are separated from one another by segments exposed to the gas environment. The YSZ is fine-grained and equiaxed as the TGO is approached. The TGO is about 100nm in thickness in the as-processed condition. The interface between the bond coat-tgo and TBC is fairly smooth but does exhibit some mild undulations. Cross-section transmission electron microscopy (XTEM) is being used to characterize the various layers in more detail. Figure 9 presents XTEM micrographs of TGO and and the inner portion of the TBC. The TGO was found to be a continuous layer of -alumina 100nm thick. The inner portion of the TBC was found to consist of fine, equiaxed grains. This zone transitioned into the typical columnar structure with distance away from the TGO. Similar morphologies have been reported for an EBPVD TBC on a diffusion aluminide bond coat with an IN 738 substrate [12]. The growth of the TGO has been found to be faster underneath the TBC than on the bond coat-only side of the specimens [13]. Figure 10 is a micrograph showing more detail of the TGO grown under the TBC for 10 hours at 1200C. The outer zone was 4

5 found to be a fine grained mixture of alumina and zirconia plus yttria. This zone has been examined in more detail using cross-section transmission electron microscopy (XTEM). Figure 11 shows the interface between the intermixed layer and the TBC. The intermixed layer consists of a continuous matrix of alumina containing spherodized paticles of YSZ. The alumina is apparently growing into the YSZ layer, breaking it down, and spherodizing the resultant particles. Figure 12 presents a schematic diagram of how this may occur assuming that the outward growth component of the alumina is, somehow, accelerated in the presence of the TBC. This process is not fully understood and it remains to be shown if it occurs for other bond coats and over what temperature range it may occur. One possibility for the rapid outward growth of the alumina layer is the presence of transient aluminas such as or [14]. However, only -alumina has been detected to this point. Changes in Bond Coat The as-processed bond coat is shown in cross-section in Figure 13. The coating thickness is 55 m.the coating consists of an outer zone of columnar grains which extend completely through the zone and have a width of approximately 20m. The inner zone of the bond coat does not have a well defined grain structure. This inner zone probably developed as nickel diffused out to the outer zone. The light phase in this inner zone is believed to be refractory metal precipitates that form as a result of the nickel removal from this zone. Compositional profiles from the surface of this coating are presented in Figure 14. These profiles indicate that the outer zone of the bond coat is -NiAl containing Pt in solution. It is important to emphasize that TBCs fabricated using the same techniques (i.e. EBPVD-YSZ, platinum-modified diffusion aluminide bond coat) but by different sources may have different microstructures. Hence, the structures presented here are typical of one source. Furthermore, even in the same source there can be some variation of structure. It is necessary to document the extent of microstructural variations and the importance of such variations to TBC lives. The microstructure of the bond coat after 125 hours of isothermal exposure at 1135C is shown in Figure 15. The depletion of Al has resulted in the nucleation of at the grain boundaries (bright areas). The bond coat after 500 hours cyclic exposure at 1135C is shown in Figure 16. The Al depletion has progressed to the point that the coating has transformed to primarily with some -phase just beginning to develop. The grain size of the is about m which is much smaller than that of the initial - phase, Figure 13. When failure occurs, especially after rather high temperature exposures, a metallic phase enriched in tantalum is evident along the fracture surface, Figure 17. This phase may be -phase, which is evident in Figure 16 just beneath the TGO. Spinel Formation Figure 18 shows an example of spinel forming above the alumina layer in the TGO. A number of investigators have reported degradation of thermal barrier coatings during prolonged exposures by the formation of spinels, apparently because of depletion 5

6 of Al from the bond coat [16, 17] or as the result of transient oxidation e.g. by oxidation of -phase in a NiCoCrAlY bond coat [18]. This section contains a thermodynamic assessment of this process for the case of oxidation of binary Ni-Al alloys (bond coats). Figure 19 is a thermodynamic stability diagram for the Ni-Al-O system at 1000C. Boundary (1) is for the Al 2 O 3 -alloy equilibrium: 2Al + 3O = Al 2 O 3 (1) boundary (2) is for NiAl 2 O 4 -alloy equilibrium: Ni + 2Al + 4O = NiAl 2 O 4 (2) and boundary (3) is for NiO-alloy equilibrium: Ni + O = NiO (3) (Note that log a Ni increases as log a Al decreases.) Reaction (4) is a three-phase equilibrium: 4Al 2 O 3 + 3Ni = 3NiAl 2 O 4 + 2Al (4) Reaction 5 is also a three phase equilibrium and is given by: NiAl 2 O 4 + 3Ni = 4NiO + 2Al (5) The equilibrium that dominates over most of this diagram is (1), namely the Al 2 O 3 -alloy equilibrium. Hence, in the case of aluminide (and MCrAlY) coatings the TGO is initially Al 2 O 3. As the aluminum activity in these coatings decreases, as occurs during high temperature exposure, the oxygen activity increases as described by boundary (1). it has been observed experimentally that, for nickel alloys containing aluminum concentrations less than about 3 wt%, that oxygen moves into such alloys faster than aluminum can diffuse out. Consequently, when the aluminum activity in the alloy decreases below about 10-11, discontinuous Al 2 O 3 particles should begin to develop in the coating just below the TGO-coating interface. As this condition continues the oxygen activity at the Al 2 O 3 -coating interface will continue to increase and, when the crossing point of boundaries (1) and (2) is reached, Al 2 O 3 at the interface can be converted to NiAl 2 O 4 via reaction (4). However, it is emphasized that this requires the aluminum activity to be lowered to a low value of less than In order to determine what phases may form during oxidation, one can construct reaction paths on the stability diagram. The actual path that is followed will be determined by kinetic processes. Figures 20 and 21 are schematic versions of the stability diagram in Figure 19 which illustrate possible reaction paths. These figures also present schematic diagrams of the resultant oxidation morphologies. Figure 20 shows a reaction path which would be typical of the case where Al depletion has not proceeded to the 6

7 extent that other oxides can form. The oxygen activity increases as one proceeds from the unaffected alloy through the Al-depleted zone into the Al 2 O 3, and aluminum activity decreases. The mechanism in this case is aluminum diffusing from the alloy to the alloy- Al 2 O 3 interface. Oxygen, or aluminum, or both are diffusing through the Al 2 O 3. The amount of nickel in the Al 2 O 3 is very small even though its activity at the alloy- Al 2 O 3 interface is high. In Figure 21 conditions are similar to Figure 20, but the alloy has been depleted of aluminum to the extent, described above, that oxygen diffuses into the alloy from the Al 2 O 3 -alloy interface and forms discontinuous oxide particles. The oxygen is supplied to the alloy by the reverse of reaction (1). The aluminum that is produced by this decomposition is oxidized by oxygen coming through the oxide, or it diffuses to the Al 2 O 3 -gas interface and is oxidized. Additionally, if nickel can diffuse through the Al 2 O 3, it is possible for NiAl 2 O 4 to form at the Al 2 O 3 -gas interface. The NiAl 2 O 4 can then continue to grow by transport through the NiAl 2 O 4 layer. This reaction path seems to describe the development of the morphology in Figure 18. Effect of Bond Coat Processing The manner in which a given bond coat is processed can also affect TBC performance. An example of this is given in Figures 22 and 23. Figure 22 presents the surface of a Pt aluminide bond coat. The as-coated surface consists of ridges that form at grain boundaries in the coating during aluminizing. The coating at right has had the ridges removed by grit blasting. The dark particles are embedded alumina grit. Figure 23 presents cross-sections of these coatings after isothermal oxidation in air for 100 hours. The as-aluminized coating has undergone preferential grain boundary oxidation which is believed to be associated with impurities in the grain boundaries, particularly refractory metals coming from the substrate during aluminizing. The grit blasted coating shows no evidence of grain boundary oxidation. Apparently, the deformation of the surface has accelerated aluminum transport and/or dispersed the impurities to the point where a continuous alumina film can form over the boundaries. Work by Gell et al. [19] has identified the oxide ridges and grain boundary oxidation as one source for the initiation of TBC failure. Effect of Exposure Atmosphere The effect of water vapor can significantly affect the time to failure, as mentioned in the discussion of Figure 6 where at 1100C the failure time in air plus 0.1 atm water vapor was approximately half that in dry air. Figure 24 shows macroscopic photographs of two specimens failed in the wet gas. The failures are edge-initiated buckles as in the case of the failures in dry gas. However, the fracture surfaces are quite different. Figure 25 shows the surface of one of the specimens after spallation. There are two significant features of this surface. First, there is a significant amount of spinel which has formed above the alumina TGO, only a trace of which was observed for the failure in dry air, Fig. 4. Second, there is virtually no spalling to the bare bond coat, which again is very different from the failure in dry air, Figs. 4 and 5. The presence of the water vapor has 7

8 resulted in substantial outward transport of nickel through the alumina to form spinel by a reaction such as that in Figure 21. Furthermore, the presence of the spinel has resulted in the fracture path occurring exclusively in oxide phases. This result is of practical significance because virtually all applications of TBCs are in combustion environments which contain water vapor as one of the combustion products. Potentials for Improved TBC Systems Identification of schemes for improving the durability of TBC systems is a difficult task. Existing state-of-the-art systems are already well engineered. For example, if it is assumed that the data in Figure 6 can be extrapolated to lower temperatures, one finds that failure times would be expected to exceed 5,000 hours at 1050C and 50,000 hours at 1000C. These temperatures are more realistic, particularly for the land-based gas turbines. Therefore, a first step in developing improved systems is to completely understand the factors which control the behavior of current systems. This avenue is being taken in a number of laboratories but the work is not yet complete. Nevertheless, the understanding that has evolved to this point suggests several directions which may result in improved systems. These are briefly described in the following. Improved Ceramic Topcoats Some approaches to improving the ceramic topcoat, such as developing nanostructured coatings to lower thermal conductivity, have already proven to be unproductive. However, the development of coatings which are more resistant to sintering would improve durability. The effect of sintering of EP-PVD and APS YSZ coatings has been illustrated in Figures 7 and 8. This process increases the effective modulus of the coating and the amount of elastic energy stored in the coating both of which can result in premature coating failure. Figure 26 shows sintering data published by Zhu and Miller [8] which compares the sintering behavior of of yttria-stabilized zirconia and yttria-stabilized hafnia. The sintering rate of the HfO 2 -based ceramic is clearly slower than that for the YSZ over the temperature range likely to be encountered in the topcoat, C. Thus, one avenue which could be explored is replacing all, or part, of the zirconia with hafnia. Modified Bond Coats The thermal expansion mismatch between the substrate/bond coat and the TGO is a major contributor to coating spallation. Recent results [20] have indicated that the incorporation of fine Al 2 O 3 and Cr 23 C 6 particles in MCrAlY coatings, produced by the detonation gun technique, resulted in substantially improved cyclic oxidation resistance, as measured by the amount of Al-depletion from the outer portion of the coating. The mechanism for this effect has not been ascertained but may well result, in part, from a better match between the thermal expansion coefficients of the coating and the protective alumina layer and the higher strength of the particle-dispersed coating which may allow it to better resist rumpling under the influence of thermal and growth stresses. While for thick substrates the thermal expansion coefficient of the bond coat may not have a strong influence on spalling, further attempts at lowering the thermal expansion coefficient of 8

9 the bond coat, while maintaining its oxidation resistance, may prove fruitful in cases where the substrate is thin. Diffusion Barriers Interdiffusion between the superalloy substrate and the bond coat results in Al depletion from the bond coat (e. g. Fig. 15 and 16), which can lead to spinel formation, and the incorporation of substrate elements into the bond coat (e. g. Fig.17), which can lead to accelerated oxidation. Indeed, the same bond coat has been observed to behave differently on different substrates [21] and the deposition of a TBC onto bulk Zr-doped NiAl resulted in much longer times to failure than are observed for typical systems [22]. A key factor in the latter observation was the absence of any interdiffusion effects. Therefore, the use of a diffusion barrier between the superalloy substrate and the bond coat has the potential for improving coating life. Nesbitt and Lei [23] have analyzed the potential benefits of this approach for a NiCoCrAlY coating and, using a finite-difference diffusion model, predict that an optimized diffusion barrier would double the coating life, based on Al depletion. This analysis cannot be directly translated to TBC systems, where the first spallation event is critical, but does suggest some potential for diffusion barriers. Diffusion aluminide bond coats are, by their nature, not amneable to the diffusion barrier approach. No Bond Coat Systems The bond coat grain boundaries can result in accelerated oxidation (Fig. 23) and have been associated with failure initiation [19]. One remedy for this situation is to deposit the TBC directly onto an oxidation resistant single crystal superalloy. Smialek [24] has compared the behavior of an APS TBC on low sulfur PWA 1484 and the same alloy with a NiCrAlY bond coat. The bond coat system failed after 1100 hours and the failure times for the no-bond coat system ranged from 500 to 1000 hours. Significant amounts of spinel were observed on the fracture surface and there was essentially no spalling to the bare substrate. The spinel formation, which is almost certainly the cause of the shorter failure times for the no-bond coat system, is likely the result of transient oxidation of the superalloy. This phenomenon, which is typical of the superalloys, is illustrated in Figure 27 which shows a layer of Ni- and Ta-containing oxides which formed during the initial exposure of the alloy and now sit on top of the continuous alumina layer. Prevention of this transient oxide layer is necessary to avoid premature spalling of a TBC deposited directly onto the superalloy. The results of a recent investigation [25] indicate that preoxidation of the superalloy in a low oxygen pressure atmosphere can eliminate the transient oxides but the results were not reproducible. However, when a thin layer of platinum was applied to the superalloy the transient oxides could be eliminated. Figure 28 presents the surface and cross-section of a Ni-base single crystal superalloy which was coated with a 7 m thick Pt layer and then oxidized in air at 1100C for 23 hours. The only oxide formed was alumina. However, particles of Pt were left on top of the alumina scale. Figure 29 shows the results for an identical specimen which was annealed in a low oxygen pressure atmosphere for 2 hours prior to the air oxidation. Apparently the initial annealing treatment allowed enough diffusion of the Pt 9

10 into the substrate that there was no Pt remaining on top of the alumina scale. Such an alloy could well be effective in a no-bond coat system. SUMMARY AND CONCLUSIONS Yttria-stabilized Zirconia (YSZ) coatings deposited by electron beam physical vapor deposition on platinum aluminide bond coats on the superalloy N5 have been oxidized at temperatures between 1100 and 1200C in air. The thermally grown oxide (TGO) that develops between the bond coat and the TBC during oxidation, as well as the bond coat and the TBC adjacent to the TGO, have been examined for both as-processed and exposed coatings using optical metallography, SEM, and XTEM. The as-processed TBC is composed of large columnar grains of YSZ that are separated from one another by segments exposed to the gas environment. The YSZ is fine-grained and equiaxed as the TGO is approached. The TGO is a 100nm thick layer comprised of columnar -alumina grains. The bond coat is 55 m thick and consists of an outer zone of columnar grains of -NiAl (with Pt in solution), which extend completely through the zone and have a width of aproximately 20m, and an inner diffusion zone. The growth of the TGO has been found to be faster underneath the TBC than on bond coats which did not have a YSZ overlayer during high temperature exposure. The TGO under the TBC develops an outer zone which is a fine-grained, two-phase mixture of alumina and YSZ. Sintering of the TBC has also been observed during high temperature exposure. This process is significant in that it can result in increases in the TBC stiffness and in the residual stress in the TBC. The bond coat transformed, during exposure, to primarily with some -phase just beginning to develop at longer times. The grain size of the is about m which is much smaller than that of the initial -phase. When failure occurred a metallic phase enriched in tantalum was evident along the fracture surface. In some instances spinel was observed to form between the TGO and YSZ topcoat as the result of the Al depletion from the bond coat. This was particularly evident for specimens exposed in atmospheres containing water vapor. Possible approaches for the development of improved TBC systems have been outlined. ACKNOWLEDGEMENT The authors gratefully acknowledge the financial support of this work by the U. S. Department of Energy through AGTSR Subcontract No SR046, the U.S Airforce Office of Scientific Reasearch through AFOSR Grant No. F , and the U. S. Office of Naval Research through ONR Contract N

11 REFERENCES 1. R. A. Miller, Thermal Barrier Coatings for Aircraft Engines - History and Directions, Thermal Barrier Coating Workshop, NASA CP 3312, 1995, p A. Maricocchi, A. Bartz, and D. Wortman, PVD TBC Experience on GE Aircraft Engines, Thermal Barrier Coating Workshop, NASA CP 3312, 1995, p S. Bose and J. DeMasi-Marcin, Thermal Barrier Coating Experience in Gas Turbine Engines at Pratt & Whitney, Thermal Barrier Coating Workshop, NASA CP 3312, 1995, p J. T. DeMasi-Marcin and D. K. Gupta, Surf. and Coatings Tech., 68/69, 1 (1994). 5. R. Milller, J. Amer. Ceram. Soc., 67, 517 (1984). 6. W. Lih, E. Chang, C. H. Chao, and M. L. Tsai, Oxid. Metals, 38, 99 (1992). 7. W. Lih, E. Chang, B. C. Wu, and C. H. Chao, Oxid. Metals, 36, 221 (1991). 8. D. Zhu and R. A. Miller, Surf. Coatings and Tech., , 114 (1998). 9. D. Zhu and R. A. Miller, Determination of Creep Behavior of Thermal Barrier Coatings under Laser Imposed Temperature and Stress Gradients, NASA Technical Memorandum , Nov C. A. Johnson, J. A. Ruud, R. Bruce, and D. Wortman, Surf. Coatings and Tech., , 80 (1998). 11. J. T. Thornton, The Measurement of Strains within the Bulk of Aged and Assprayed Thermal Barrier Coatings Using Syncrotron Radiation paper presented at ICMCTF, San Diego, April, O. Unal, T. E. Mitchell, and A. H. Heuer, J. Amer. Ceramic Soc., 77, 984 (1994). 13. M. J. Stiger, N. M. Yanar, F. S. Pettit, and G. H. Meier, Mechanisms for the Failure of Electron Beam Physical Vapor Deposited Thermal Barrier Coatings Induced by High Temperature Oxidation, in Elevated Temperature Coatings: Science and Technology, J. M. Hampikian and N. B. Dahotre, eds., TMS, Warrendale, PA, 1999, p D. A. Clarke, V. Serge, and M-Y He, Precursor to TBC Failure Caused by Constrained Phase Transformation in the Thermally Grown Oxide, in Elevated Temperature Coatings: Science and Technology, J. M. Hampikian and N. B. Dahotre, eds., TMS, Warrendale, PA, 1999, p I. G. Wright, B. A. Pint, W. Y. Lee, K. B. Alexander, and K. Prüssner, Some Effects of Metallic Substrate Composition on Degradation of Thermal barrier Coatings in High Temperature Surface Engineering, J. Nicholls, ed., Institute of Materials, London, in press. 16. Mutasim, C. Rimlinger, and W. Brentnall, Characterization of Plasma Sprayed Electron Beam-Physical Vapor Deposited Thermal Barrier Coatings, Reprint of Paper presented at the International Gas Turbine & Aeroengine Congress & Exhibition, Orlando, FL, Amer. Soc. of Mech.Eng., June, E. A. G. Shillington and D. R. Clarke, Acta Mater., 47, 1297 (1999). 18. W. J. Quadakkers, A. K. Tyagi, D. Clemens, R. Anton, and L. Singheiser, The Significance of Bond Coat Oxidation for the Life of TBC Coatings, in Elevated 11

12 Temperature Coatings: Science and Technology, J. M. Hampikian and N. B. Dahotre, eds., TMS, Warrendale, PA, 1999, p M. Gell, K. Vaidyanathan, B. Barber, J. Cheng, and E. Jordan, Met. and Mater. Trans A, 30A, 427 (1999). 20. T. A. Taylor and J. K. Knapp, Surf. and Coatings Tech., 76-77, 34 (1995). 21. U. Kaden, C. Leyens, M. Peters, and W. A. Kaysser, Thermal Stability of an EB- PVD Thermal Barrier Coating System on a Single Crystal Nickel-base Superalloy, in Elevated Temperature Coatings: Science and Technology, J. M. Hampikian and N. B. Dahotre, eds., TMS, Warrendale, PA, 1999, p B. A. Pint, I. G. Wright, W. Y. Lee, Y. Zhang, K. Prüssner, and K. B. Alexander, Mater. Sci. and Eng.,A245, 201 (1998). 23. J. A. Nesbitt and J-F. Lei, Diffusion Barriers, to Increase the Oxidative Life of Overlay Coatings, in Elevated Temperature Coatings: Science and Technology, J. M. Hampikian and N. B. Dahotre, eds., TMS, Warrendale, PA, 1999, p J. L. Smialek, Toward Optimum Scale and TBC Adhesion on Single Crystal Superalloys, in High Temperature Corrosion and Materials Chemistry, P. Y. Hou, M. J. McNallan, R. Oltra, E. J. Opila, and D. A. Shores, eds., The Electrochem. Soc., Pennington, NJ, 1998, p M. Topping, M.S. Thesis, Univ. of Pittsburgh,

13 LIST OF FIGURES Figure 1. Schematic diagram showing the increase in metal temperature in gas turbines resulting from superalloy development and the effective increase associated with the introduction of thermal barrier coatings. Figure 2. Schematic diagram of coated specimens used in this investigation. 13

14 Figure 3. Cross-section SEM micrograph of two as-processed TBC systems; An EB- PVD topcoat with a Pt modified aluminide bond coat (top) and an APS YSZ topcoat with a NiCoCrAlY bond coat (bottom). 14

15 Figure 4. Macroscopic photograph of the failure of an EB-PVD topcoat with a Pt modified aluminide bond coat (top) and an SEM micrograph of the bond coat surface (bottom) after 1287 cycles at 1100 C in dry air. 15

16 Figure 5. Underside of the spalled TBC from the specimen in Figure 4. 16

17 Figure 6. The effect of exposure temperature on the time to failure of an EB-PVD TBC with a Pt modified aluminide bond coat 17

18 Figure 7. Micrographs showing EBPVD TBC sintering at 1200C. Note the decrease in the number of channels after exposure. 18

19 Figure 8. Micrographs showing APS TBC sintering at 1200C. Note the decrease in the number of microcracks after exposure. 19

20 Figure 9. XTEM brightfield micrograph of the TGO and TGO/TBC interface in an as-processed EB-PVD TBC with a Pt modified aluminide bond coat. The TGO is a continuous layer of columnar grains indicated to be -Al 2 O 3 by SAD. The TBC in the interfacial region consists of fine grains of YSZ. 20

21 Figure 10. Cross-section SEM micrograph showing the intermixed region in the TGO adjacent to the TBC after 20 hours at 1200C. 21

22 Figure 11. XTEM bright field images of a specimen exposed for 10 hours at 1200C showing a) the TGO columnar and equiaxed grains and b) the TGO/TBC interface showing the incorporation of YSZ particles into the outward growing TGO. 22

23 Figure 12. Schematic diagram of proposed diffusion paths from the effects of the presence of the YSZ on the TGO. Figure 13. Optical photomicrograph of the Pt aluminide bond coat on the asprocessed specimens. The coating consists of an outer zone of columnar - NiAl and a relatively featureless inner zone. The dark lines in the coating that appear as cracks are believed to be artifacts created by etching. 23

24 Figure 14. Compositional profiles for Al, Ni, and Pt across the outer zone of the asprocessed bond coat Al Ni Pt at % Distance (m) Figure 15 Optical Micrograph showing the microstructure of the bond coat after 125 hours isothermal exposure at 1135C. 24

25 Figure 16 Optical Micrographs showing the microstructure of the bond coat after 500 hours cyclic exposure at 1135C. The two micrographs show the same coating but with two different etches which reverse the contrast. 25

26 Figure 17 Cross-section and fracture surface revealing metallic particles at the TGO/bond coat interface. 26

27 Figure 18. Cross-section of a bond coat showing spinel formation on top of the TGO. 27

28 Figure 19. Thermodynamic stability diagram for the Ni-Al-O system. The various boundaries indicate equilibria between nickel-aluminum alloys and different oxides. 28

29 Figure 20. Reaction path (dashed line) for an Al 2 O 3 scale growing on a nickelaluminum alloy. 29

30 Figure 21. Reaction path for a Ni-Al alloy leading to spinel formation on top of the alumina. Development of this morphology requires substantial Ni transport through the alumina layer. 30

31 Figure 22. Surface of an as-processed Pt-modified aluminide bond coat showing ridges at the coating grain boundaries (top) and a coating from which the ridges have been removed by grit blasting (bottom). 31

32 Figure 23. Cross-sections of a grit blasted (top) and as-aluminized Pt-modified aluminide bond coat after isothermal oxidation at 1100C for 100 hours (bottom). 32

33 Figure 24. Macroscopic photographs of the failures of EB-PVD TBCs with Pt modified aluminide bond coats after 652 cycles at 1100 C in air which contained 0.1 atm. water vapor. 33

34 Figure 25. Bond coat surface of one of the specimens in Figure 24 after removal of the spalled TBC. Figure 26. Shrinkage strains for yttria-stabilized zirconia and hafnia after sintering at various temperatures [8]. 34

35 Figure 27. Cross-section through a Ni-base superalloy after 24 hours oxidation at 1100C showing transient oxides on top of a continuous alumina layer. 35

36 Figure 28. Surface (a) and cross-section (b) of a Ni-base single crystal superalloy which had been coated with a 7 m thick layer of Pt and oxidized for 23 hours at 1100C in air. 36

37 Figure 29. Surface (a) and cross-section (b) of a Ni-base single crystal superalloy which had been coated with a 7 m thick layer of Pt, annealed in a low oxygen partial pressure atmosphere at 1100C for 2 hours, and oxidized for 20 hours at 1100C in air. 37