Microstructure, composition and property relationships of plasma-sprayed thermal barrier coatings

Surface and Coatings Technology, 50 (1992) 141-149

141

Microstructure, composition and property relationships of plasmasprayed thermal barrier coatings* R. T a y l o r a n d J. R. B r a n d o n t Manchester Materials Science Centre, Manchester (UK)

Paul Morrell Rolls-Royce plc, Derby (UK) (Received April 22, 1991; accepted July 19, 1991)

Abstract The benefits deriving from the use of thermal barrier coatings are briefly reviewed. Properties important for a thermal barrier coating are discussed and reasons for the selection of ZrO2-based alloys highlighted. Candidate zirconia alloys are critically evaluated in terms of their thermal and mechanical properties. Since coatings currently applied consist of a tetragonal phase, the long-term stability of the phase structure and the avoidance of the monoclinic phase is addressed, as is the effect of long-term ageing on the thermal properties. The relationship between microstructure and properties of plasma coatings is reviewed. Future trends, such as extending areas of application in jet engines, use of alternative alloys, and methods of application and optimization of plasma sprayed coatings by modelling the plasma spray process, are also briefly discussed.

1. I n t r o d u c t i o n i100

The performance and efficiency of a gas turbine engine are directly related to the operating temperature. Since its inception the drive to improve the gas turbine engine has resulted in the development of new alloys with improved properties and has led to significant advances in engine design such as blade and film cooling for the turbine section. The use of thermal barrier coatings on gas turbine components can lead to further increases in operating temperature but, more realistically, may augment cooling effectiveness. The benefits of using thermal barrier coatings on critical components in the turbine section have been reviewed by a number of authors [1-4]. Since the 1970s, coatings have been applied to combustor cans and by the 1980s they were also being used in low risk regions in the turbine section such as vane platforms. Use of such coatings offers reductions in meta! temperature of around 100 °C (Fig. 1). Thermal barrier coatings currently used consist of a ceramic insulating layer applied over an intermediate metallic bond coat layer. Both the ceramic and metallic bond *Paper presented at the 18th International Conference on Metallurgical Coatings and Thin Films, San Diego, CA, April 22-26, 1991. tpresent address: ERDC, Capenhurst, Chester, UI,L

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coat layers are usually applied by plasma spraying because this is economically capable of producing durable and reproducible coatings. The production properties and service behaviour of thermal barrier coatings are complexly interlinked and depend on many factors. Amongst these we may list plasma spray parameters, powder quality, bond coat and substrate chemistry, ceramic coating composition, microstructure, layer thicknesses and the operating thermal regime and environment. There is no flexibility concerning the choice of substrate, which has to be a nickel-based superalloy. The

© 1999.- Elsevier Sequoia. All rights reserved

142

R. Taylor et al. / Plasma-sprayed thermal barrier coatings

property of most relevance is the thermal expansion of the substrate, which dictates the thermal expansion needed from the ceramic. The bond coat is an essential part of the system because of its stress-relieving properties and its oxidation resistance, and because it provides a "key" for the firm adhesion of the ceramic coating. It is now recognized that ultimate failure of the thermal barrier coating system is strongly influenced by bond coat oxidation which is a crucial factor in life prediction models. Whilst recognizing the importance of the bond coat to the thermal barrier coating systems, a candidate material for the ceramic must possess certain criteria [5, 6]. It should be refractory, chemically inert, possess good mechanical strength and thermal shock resistance, have good wear and erosion resistance, be phase stable but most importantly of all possess a low thermal conductivity and a thermal expansion coefficient similar to that of the nickel-based superaUoy substrate. Many ceramics can fulfil the first five criteria and many ceramics have low thermal conductivities. However, there are relatively few ceramics which possess high thermal expansion coefficients (11-13x 10 -6 °C) and this has focused attention on zirconia. Although zirconia possesses a low thermal conductivity and a suitable expansion match with the substrate, the polymorphism of zirconia necessitates that it be alloyed with other oxides. Hence more candidate materials become available depending on the alloy systems chosen. Although it is recognized that the ultimate tests of confidence in a thermal barrier system are rig and engine running, a great deal of information about the suitability of candidate systems can be obtained by analysis of the physical, structural and thermal properties of materials prior to extensive and expensive test programmes. The evaluation of the mechanical and thermal properties of materials from two alloy systems ZrO2-Y203 and ZrO2-CeO2, the effect of temperature on phase stability and its consequent effect on the important physical properties, and the implications for the service life of coatings in modern jet engines will be discussed in this paper.

sprayed condition this single cubic phase is basically a metastable system. When coatings of the magnesiastabilized zirconia are thermally cycled between 20 and 1200 °C, the thermal diffusivity progressively increases with each cycle [7] (Fig. 2). Similar effects have been noted for ZrO2--CaO alloys [8-10]. Brandt and Neuer [8] noted a factor of two increase in thermal diffusivity for a ZrO2-CaO coating isothermally annealed at 1300 °C for 2 h. These increases are due to the precipitation, from the solid solution, of the stabilizing oxide. Precipitation has been noted at triple points and in the vicinity of intersplat boundaries in ZrO2--MgO alloys [7]. The removal of MgO from solid solution in ZrO2 will increase the thermal conductivity in two ways; (a) a two-phase ceramic will always exhibit a higher thermal conductivity than a solid solution, especially when one phase (MgO) has a higher thermal conductivity; (b) removal of MgO from the solid solution reduces point defect scattering of phonons in the zirconia phase. A further factor is that the volume change associated with the precipitation of MgO or CaO is likely to produce sub-critical microcracking in the vicinity of the precipitate. Some degree of strain accommodation in plasma-sprayed systems is due to the porosity and microcracking present as a result of the quenching of individual splat particles during plasma spraying [11]. Hence the enhancement of this crack network and the consequent improvement in strain tolerance is the major reason why magnesia-stabilized zirconia remains a viable coating which survives when sprayed onto static components in a jet engine. On the debit side is the fact that the thermal conductivity increases; each thermal cycle reduces the effectiveness of the coating as a thermal barrier. Whilst the coating still gives some protection to the substrate, the conductivity increase does to a large degree nullify the primary reason for its existence. v 6 m

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2. Zirconia alloys

Early success was achieved using plasma-sprayed coatings, of magnesia-stabilized zirconia (MSZ) containing 25 wt.% MgO which consisted of a 100% cubic phase. Such coatings have been successfully used since the early 1970s on aero-engine components. However, the various equilibrium diagrams reported for the ZrO2-MgO system show that the equilibrium phases are, at temperatures less than 1400 °C, either monoclinic or tetragonal zireonia plus MgO. Hence in the as-

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IZ Taylor et aL / Plasma-sprayed thermal barrier coatings

Alternative systems need to be considered and the two that have received most attention are ZrO2-Y203 and ZrOE-CeOz. Zirconia fully stabilized with yttria shows no signs of hysteresis in its thermal transport behaviour and showed no signs of the precipitation of Zr3Y4012 from solid solution as predicted by Stubican et al. [12] even when aged for 100 h at 1600 °C [13]. However, burner test lives [14] or thermal cycling test lives [1], showed this material to be markedly inferior in performance to calcia- or magnesia-stabilized zirconia. Hence detailed investigations have been carried out on two alloy systems to determine the optimum compositions and characterize their long-term durability [15-18]: (a) zirconia partially stabilized with yttria, compositions in the range 6-12 wt.% Y203; (b) zirconia alloyed with CeO2, compositions in the range 12-25 wt.%

to Miller et al. [13] to resolve the relative proportions of tetragonal and cubic phases. For monoclinic-tetragonal mixtures the integrated intensity ratio of the monoclinic phase Xm is defined by Im(lll) +lm(ll~.) Xrn= l m ( l l l ) + I m ( l l i ) + I t , c(111) to yield the volume fraction

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It should, at this point, be noted that because of the success achieved with binary zirconia systems, little attention has been paid to alternatives. However, ternary systems such as ZrOE-YEO3-CeO2, ZrOE-YEO3-MgO, and even non ZrO2 coatings, may still be worthy of investigation.

3. Experimental details

Powders used in these investigations were fabricated by a co-precipitation process by Magnesium Electron Ltd. The thermal barrier coatings were plasma sprayed using a Metco 7 MB gun to a set of standard spray parameters prescribed by Rolls-Royce. For thermal diffusivity measurements, coatings 150-300 /zm thick were plasma sprayed onto nickel-based superalloy substrates. Thermal diffusivity measurements were made using the laser flash technique [19]. Specimens for thermal expansion were prepared by spraying a coating approximately 1 mm thick onto a nickel alloy substrate which was then removed by etching. Thermal expansion measurements were made in an alumina pushrod dilatometer [16] at a constant ramping rate of 2 °C min- 1. Phase compositions of the coatings were measured in the as-sprayed state or, subsequent to testing, by Xray diffraction using a Philips horizontal diffractometer PW1380 using nickel filter Cu Ka radiation. Three scans were used: (1) a continuous scan 2 0 = 10°-80 ° at 0.1 s-l; (2) a slow scan at 0.005 ° s -1 of the region 20=71°--76 ° to examine (400) type reflections; (3) a slow scan at 0.01 ° s -1 of the region 20=27.32 ° to examine (11i) and (111) monoclinic and (111) tetragonal reflections. Peak deconvolution was done manually using an iterative method. Phase analysis [17, 18] was carried out using the analysis of Toraya et al. [20] to determine tetragonal and monoclinic proportions, and that due

143

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Two types of mechanical property test were used: tensile adhesion and fracture toughness. In the tensile test a bar was plasma sprayed with bond coat and ceramic. A second bar was plasma sprayed only with a bond coat. These were glued together on a special jig using an epoxy resin having a tensile strength when fully cured of greater than 30 MPa. These composite bars were tested in tension in an Instron testing machine. In the fracture toughness test, the critical strain energy release rate (for either adhesive or cohesive failure) was measured using the constant bending moment test devised by Friedman et al. [21] and developed by Becher and Newall [22]. The two bars, one containing bond coat and ceramic, the other containing only bond coat, were joined using the same adhesive. 4. Results and discussion 4.1. Phase analysis Phase analysis data for the as-sprayed coatings are listed in Table 1. In the case of the yttria series of alloys, compositions of 8-12 wt.% Y203 consisted wholly of the non-transformable tetragonal t' phase [13, 23]. Whereas the 6 wt.% yttria alloy contained only a trace T A B L E 1. Phase analysis data for as-sprayed coatings Alloy 4 6 8 10 12 12 15 20 25

wt.% wt.% wt.% wt.% wt.% wt.% wt.% wt.% wt.%

Y203 Y203 Y203 Y203 Y2Oa CeO 2

CeO2 CeO 2 CeO2

Vt ( % )

Vm

Vc

a(nm)

c(nm)

c/a

58 99 100 100 100 56 72 90 87

42 1 44 28 10 -

-13

0.5102 0.5104 0.5113 0.5118 0.5120 0.5114 0.5110 0.5124 0.5137

0.5174 0.5174 0.5160 0.5154 0.5146 0.520 0.520 0.520 0.5209

1.014 1.0126 1.010 1.007 1.005 1.017 1.017 1.017 1.014

144

R. Taylor et al. / Plasma-sprayed thermal barrier coatings

of monoclinic phase, the 4 wt.% yttria coating contained a much larger monoclinic fraction. The c/a ratio decreases with increasing yttria content, as noted by Scott [23] and Miller et al. [13]. Using data for the measured c and a axis lattice parameters [24] extrapolated to room temperature, the relationship deduced is [17] wt.%'Y203

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This is similar to the relationship deduced by Miller et al. [13] and these are plotted in Fig. 3. Quantitatively, these X-ray data are in accord with the predictions of Anderson et al. [25] that two martensitic transformations exist in this system Ms t ~ m and Ms c - f - ~ t . In contrast, for the zirconia-ceria system the c/a ratio for the tetragonal phase does not change for compositions up to 20 wt.% CeO2 which exist as a mixture of discrete tetragonal and monoclinic phases. The most widely accepted equilibrium diagram is due to Tani et al. [26] which predicts that the equilibrium phases should be cubic and monoclinic. The absence of any cubic phase suggests that a similar martensitic transformation Ms c ~ t occurs in this system, as well as the familiar Ms t ~ m transformation. The 25 wt.% CeO 2 alloy existed as a mixture of cubic and tetragonal phases. 4.2. Thermal diffusivity Thermal diffusivity data for the wholly tetragonal yttria alloy compositions show a systematic decrease with increasing yttria concentration (Fig. 4). However, the diffusivity of the 6 wt.% Y203 alloy is similar to that of the 10 wt.% Y203 alloy. All alloys exhibited values that were roughly independent of temperature. Values lay within the range 2.5-4.5x10 -7 m E s -1 (equivalent to a thermal conductivity of 0.6-1.1 W m - ~ K-l). In an earlier study [16] of the effect of plasma spraying different powder particle sizes and morphologies, similar values for diffusivity were noted. Two effects can influence the thermal diffusivity or conductivity of a thermal barrier coating. The first is

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the phase structure and the second is the crack morphology. A survey of thermal properties of dense polycrystalline ceramics [27] demonstrates that for any given alloy system, monoclinic zirconia has a higher thermal conductivity than either cubic or tetragonal zirconia. Hence the anomalously low result for the 6 wt.% Y203 alloy cannot be explained by the monoclinic content. Thermal conductivity and diffusivity are structuresensitive properties, and changes in porosity or crack morphology could mask any changes due to a composition difference. The effect of microcracks on thermal conductivity has been analysed by a number of authors [28-30]. Hasselman and Singh [29] derived relationships to predict the effect of microcracks on the thermal conductivity. For cracks preferentially oriented with their long axis normal to the heat flow direction, the effective thermal conductivity iteer is given by ;re, = X(1 + 8Nb 3/3)- 1 where it is the bulk thermal conductivity, N the number of cracks per unit volume and b the radius of revolution of the "penny"-shaped cracks. In general, thermal conductivities of plasma-sprayed coating are one third to one fifth the value for similar material in the form of sintered ceramic [27]. However, changes in crack morphology could mask any change due to compositional or phase differences and this probably explains the low value for the ZrO2+6 wt.% Y203 alloy. In contrast, the zirconia--ceria alloys show a systematic trend in that the diffusivity decreases with increasing ceria content from 3.0 to 1.5x10 -7 m 2 s -1 (Fig. 5) (equivalent to a thermal conductivity of 0.5-1.0 W m K-l). It is also noteworthy that these coatings all exhibited hysteresis in their thermal diffusivity during heating and cooling. 4.3. Thermal expansion If, however, we consider the thermal expansion behaviour of the yttria partially stabilized coating, then a systematic trend is evident (Fig. 6). Expansion curves

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Fig. 5. Thermal diffusivityof ZrO2--CeO2 coating. clinic---, tetragonal (m---,t) on heating; (c) tetragonal ~ monoclinic (t ~ m) on cooling. For the 12-20 wt.% CeO2 alloys there is a significant increase in the monoclinic content which results in substantial volume changes resulting in length increases of 0.15% (20 wt.% CeO2) to 0.9% (12 wt.% CeO2).

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The results for the mechanical property tests on ZrO2-Y203 alloys and those for the ZrO2-CeO2 alloys are listed in Table 2. If Young's modulus is known, a fracture toughness value may be calculated from the measured critical strain energy release rate. Assuming a Young's modulus for the coating of 25.7 GPa [31, 32], the fracture toughness of ZrO2-Y203 alloys is roughly constant at approximately 1.9-1.97 MPa m -lt2 for the 6 and 8 wt.% alloys but decreases with increasing Y203 concentration. The tensile strength data show a systematic decrease with increasing yttria content. Examination of the fracture surfaces of the specimens showed that fracture occurred either within the ceramic or at the bond-coat-ceramic interface, i.e. mixed adhesive--cohesive failure. The fracture toughness data for zirconia-ceria alloys showed much higher values ranging from 3.6-4.6 MPa m -1/2. However, only for the ZRO2+25 wt.% CeO2 alloy did failure occur either within the ceramic or at the ceramic-bond-coat interface. For the other alloys failure predominantly occurred either at the glue-ceramic interface or at the bond-coat-glue interface. This suggests that the actual values are somewhat higher than this and that a stronger adhesive is needed. However, in tensile strength tests for the ZrO2--CeO2 alloys, failure was through the ceramic coating. Values for the tetragonal-monoclinic 12-20 wt.% CeO2 alloys showed values ranging from 12.7 to 17.9 MN m -2 but for the 25 wt.% CeO2 alloy a very low value of 5.6 MN m -2 was recorded.

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for the 8, 10, 12 wt.% Y203 alloys which consist of 100% t' phase are virtually identical at a value of approximately 11 x l0 -6 °C-1. For the 6 wt.% Y203 alloy a contraction in length occurred above 250 °C which is attributed to the transformation of the monoclinic to the tetragonal phase. The monoclinic phase was not subsequently detected, the reverse transformation tetragonal to monoclinic did not occur on cooling and this resulted in a length decrease of 0.3%. The 4 wt.% alloy exhibited two transformations; at around 200--250 °C the retained transformable tetragonal phase transforms to the monoclinic phase and transforms back to the tetragonal phase above 350 °C. From 800 to 1000 °C, a linear expansion coefficient of 11 x 10-6 was noted. However, on cooling, linear expansion was noted down to 400 °C. Below 400 °C, where the tetragonal-monoclinic transformation occurred, there was a large expansion which resulted in a permanent increase in length of 0.8%. This is consistent with an observed increase in monoclinic content from 42% to 66%. However, for the ZrO2-CeO2 alloys the thermal expansion data (Fig. 7) indicate that only the 25 wt.% CeO2 coating, which has an expansion coefficient of 10.5-12.0 × 10 -6 K-1, exhibits linear behaviour. All the other coatings exhibit non-linear behaviour attributable to the following transformations: (a) retained tetragonal--* monoclinic (t ~ m) on heating; (b) mono-

4.5. Microstructure and the plasma spray process

A plasma-sprayed coating is built up layer by layer of individual splat particles which have melted or become

R. Taylor et al. / Plasma-sprayed thermal barrier coatings

146

TABLE 2. Tensile strength and fracture toughness data for ZrO2-Y203 and ZrO2-CeO z alloys Alloy

6 8 10 12 12 15 20 25

wt.% wt.% wt.% wt.% wt.% wt.% wt.% wt.%

Y203 Y203 Y203 Y203 CeO2 CeO2 CeO2 CeO2

Tensile strength (MPa m -z)

Standard deviation o- (MPa m -2)

Fracture toughness (MPa m -If2)

Standard deviation o (MPa m -xtz)

14.0 10.0 7.0 5.2 17.9 12.7 14.2 5.6

3.2 4.1 2.6 1.9

1.90 1.97 1.61 1.19 3.79 4.03 4.62 3.6

0.24 0.18 0.15 0.18 0.3 0.38 0.29 0.45

semi-molten during passage through the plasma arc. The microstructure, thermal conductivity, cohesive and adhesive strength, and modulus of a coating will be dependent on the bonding between these splats or particles and their packing arrangement and density. The two principal parameters which govern this are the velocity of the particles on impact and the viscosity of the particles as they flow. These are not controlled directly in the plasma spray process but are influenced by the temperature of the particle and the velocity of the arc. These are in turn influenced by the parameters of the process that can be controlled: the plasma gas flow and current density, the torch geometry, powder morphology, density and Biot number, sprayingvariables such as deposition temperature of the substrate, surface traverse rates, thickness per pass etc. A number of attempts are being made to model the plasma spray process [33-35] and a number of possible particle deposit types have been identified [36]. These arise owing to the fact that different powder particles follow different trajectories through the plasma torch before being quenched on impact with the substrate. Segmentation microcracking and internal stresses are also generated during quench cooling of individual splat particles. This gives rise to a microstructure which contains a high density of cracks parallel to the substrate (Fig. 8), a network of quench cooling cracks normal to the substrate and some porosity. Coating densities were approximately 90% of theoretical density. The presence of the pores, microcracks and interlamellar cracks has been shown to reduce the elastic modulus of plasma-sprayed A1203 to around 20% of the value of a monoclinic ceramic, and to reduce dramatically Poisson's ratio to less than 0.1 [37]. This microstructure gives the coating a degree of flexibility, enabling it to withstand the strains due to thermal expansion mismatch and thermal cycling. Even though it is generally agreed that the life-limiting factor in predicting thermal barrier coating lives is oxidation of the bond coat [38, 39], it is important that the strain accommodation capability of the ceramic is not impaired

Fig. 8. Microstructure of a typical plasma-sprayed coating.

either during thermal cycling or by high temperature exposure. There are likely to be two possible causes for this: sintering and microcrack healing, or a significant increase in microcracking which would weaken the coating, causing it to disintegrate. In fact, the available evidence suggests that for ZrO2-MgO alloys, the converse may be true, i.e. the precipitation of MgO produces a controlled degree of random microcracking that enhances the properties. Accordingly, it is important to evaluate the effect on coating properties of long-term high temperature exposure for which in the first instance we need to consider the phase stability of these alloys. 4.6. Annealing experiments Samples of all materials tested were annealed for periods of up to 100 h at temperatures in the range 1200-1600 °C [17, 18]. When samples of t' phase are annealed at 1200-1600 °C, the original t' phase should, in theory, decompose to form two equilibrium tetragonal and cubic phases. On cooling to room temperature, the high yttria phase may be retained as cubic or it may transform to a high yttria t' phase; the low yttria phase may transform to the monoclinic phase. A fairly complex series of time- and temperature-dependent transformations were observed which support the

R. Taylor et al. / Plasma-sprayed thermal barrier coatings

hypothesis that low yttria t' (tt') and high yttria t' (t2') phases are formed [17]. A typical example showing the decomposition of ZRO2+8 wt.% ½Oz at 1500 °C as a function of annealing time is shown in Fig. 9. Only for the 4 and 6 wt.% yttria coatings was any monoclinic phase observed. However, it was felt that the label non-transformable may not be strictly applicable to the low yttria phase after certain annealing treatments, and that rapid cooling rates during air cooling may have suppressed transformation of the low yttria tt' phase formed during annealing. This was, in fact, found to be the case. Figure 10 shows the thermal expansion behaviour of a specimen annealed at 1400 °C for 100 h. The following may be noted: (a) retained tetragonal ~ monoclinic t ~ m at 300 °C on heating; (b) monoclinic-+ tetragonal m ~ t at 600-750 °C on heating; (c) tetragonal-->monoclinic t--+m at 400--100 °C on cooling; (d) the tl' phase has disappeared following the expansion. (Fig. 11). These observations are also substantiated by the work of Van Valzah and Eaton

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[40] who showed that the monoclinic content is dependent on cooling rate. Similar annealing studies followed by thermal expansion measurements were carried out on the ZrO2-CeO2 series of alloys [18]. Since the alloys containing less than or equal to 20 wt.% Ce02 were shown to contain significant amounts of monoclinic phase, only the 25 wt.% CeO2 alloy, which as-sprayed contained a wholly tetragonal phase, is of especial interest. Only after annealing for 100 h at 1600 °C was a significant (13%) monoclinic content detected. This produced a slight non-linearity in the expansion behaviour at low temperatures but was insufficient to produce a net dimensional change.

4.7. In service experience Zirconia partially stabilized with 8wt.% ½03 (8% PYSZ) was chosen to replace the older MSZ systems in service based on the stability of its thermal properties and its superior mechanical behaviour. These coatings have now been flying in RB2 eleven-E4 engines since 1981 and have out performed the MSZ systems in controlled-service introduction trials [41] and have now amassed up to 15 000 h in service (Fig. 12). The optimum performance of the plasma-sprayed coatings has been achieved by the 8% PYSZ coating where the bond coat has been deposited by vacuum plasma spraying (VPS). This is because the life of the thermal barrier coating system is primarily a function of the oxidation kinetics of the bond coat [38]. These engine trials, however, operate the ceramic up to temperatures in the region of 1200 °C and therefore it would not be expected for there to be any marked change in the crystallographic structures, as the metastable t' phase does not transform to the high and low YzO3 forms, owing to the sluggish rate of

148

R. Taylor et al. / Plasma-sprayed thermal barrier coatings

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TIME, hours

Fig. 12. Rolls-Royce experience of various thermal barrier coating systems.

14 15

16

transformation, at these temperatures. However, if these ceramics were to operate at temperatures significantly higher than 1200 °C, the phase changes noted in this paper would be expected to occur. The instability exhibited by the systems of ZrO2 and 12-20 wt.% CeO2 has meant that they could never be realistically considered as thermal barrier coating materials for gas turbines. The 25 wt.% CeO2 material offered some benefits as its thermal and mechanical properties were reasonable, but independent erosion tests showed that its erosion resistance was significantly inferior to that of the PYSZ systems [39]. However, the ZRO2-25 wt.% CeO2 coating did offer high temperature phase stability, even following ageing at 1600 °C for 100 h [18]. This may be a possible alternative to the 8% PYSZ system for high temperature applications should the formation of a high and low Y 2 0 3 t' phase and subsequent transformation of the low Y203 t' phase to the monoclinic phase cause significant cracking in the coating, sufficient to cause spallation.

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