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State of the art on high-temperature corrosion-resistant coatings
Materials Science and Engineering, A 120 (1989) 13-24
13
State of the Art on High-temperature Corrosion-resistant Coatings R. M]~VREL ONERA, Materials Science Department, BP 72, 92322Chtitillon (France) (Received July 10, 1989)
Abstract
A high-temperature protective coating must meet several criteria: provide adequate environmental resistance, be chemically and mechanically compatible with the substrate, be applicable. Comprehensive reviews on high-temperature coatings have appeared regularly since the early 1970s. Our purpose is not to recapitulate the material covered therein but rather to focus on recent trends, and point out some research perspectives. Historically the development of high-temperature protective coatings has been linked with the evolution of demanding applications such as superalloy components in gas turbine engines; the searches for better performance (higher inlet temperatures, longer lifetimes, etc.) and for cost-saving solutions (use of contaminated low-grade fuels) have been the main incentives for developing the different coatings now available in production: simple and "modified" diffusion coatings, overlay coatings, thermal barriers. In recent years research and development activity has been concentrating on the following points. (a) Degradation mechanisms in high-temperature corrosion of metallic coatings; basic studies on the growth mechanisms of oxide scales, for exampie, are still required, in particular to understand the role of addition elements such as platinum and palladium. (b) Alternative techniques for depositing MCrAIY coatings; electrolytic codeposition and electrophoresis, for example, have been developed at the laboratory stage and these permit the deposition of MCrAIY coatings with claimed economic and technical advantages over processes already in production. (c) Thermal barrier coatings; ceramic coatings have been applied to sheet metal combustor components for about 15 years; only recently have they been used in the turbine section. Two challenges remain though to exploit these coatings on turbine blades: improve their reliability and, in the case of 0921-5093/89/$3.50
stationary gas turbines, their hot corrosion resistance. Both structural and mechanical approaches are required to determine, in particular, the role of microstructure, microcracking, porosity, residual stresses, oxidation of the bond layer in the degradation mechanisms of these coatings. (d) Mechanical properties of coated systems; the intrinsic mechanical properties of coating materials are still poorly described and the lack of information hinders the adequate modelling of the behaviour of coating-substrate composite systems. In parallel, an increasing activity is noted concerning the design and development of high-temperature coatings for protecting materials other than superalloys, for instance ceramic composites and titanium-based alloys.
1. Introduction
Materials for high-temperature applications are generally selected for specific properties such as strength, creep, mechanical and/or thermal fatigue. During their use at high temperatures these properties may degrade as a result of interactions with corrosive environments. One solution to avoid this alteration is to protect the component with a coating. Most high-temperature coatings rely on the formation of a protective oxide scale by interaction with the environment. The role of this scale is to isolate the coating from the aggressive environment and thereby limit the degradations due to high-temperature corrosion. Obviously this scale must fulfil several conditions: be stable, slow-growing, dense, adherent (often in cyclic conditions). In practical applications only three oxides correspond to these criteria: alumina (A1203) , chromia (Cr203) and silica (SiO2). The fact that chromia can form volatile suboxides at temperatures above about 900 °C generally restricts its use to somewhat lower temperatures, especially with high-velocity gases. Silica scales © ElsevierSequoia/Printed in The Netherlands
14 confer protection on ceramic coatings (SIC, Si3N 4 for example) deposited on carbon-based materials up to about 1800°C [1]. However, the use of silica-forming metallic alloys as coatings has been limited. This is due to the fact that in order to form a silica scale, a metallic coating must contain a relatively large amount of silicon. But silicon, which diffuses rapidly in most alloys, may form low-melting phases and/or brittle silicides by interaction with substrates. As in the case of chromia-forming coating alloys, the use of silicaforming metallic coatings must be restricted to rather low temperatures. These considerations explain why most high-temperature coatings (simple or modified aluminides, MCrA1Y overlays) rely on the formation of an alumina scale to ensure the protection of an underlying component. Comprehensive reviews on high-temperature protective coatings have regularly appeared since the early 1970s [1-6]. Our purpose is not to recapitulate the material covered therein, but rather, on the basis of published material and some personal communications, to focus on recent trends and to attempt to point out some research perspectives in the first part of this paper. The second part is devoted to substrate-coating compatibility considerations. 2. Main types of coating As mentioned above, most coatings rely on the isolation by a "barrier" of the part to be protected from the environment. The role of this barrier is to minimize the diffusion of gaseous or liquid species towards the component and conversely to prevent elemental diffusion from the alloy towards the external surface where they could react. Most generally the barrier is dynamic in the sense that it forms by interaction of the coating with the oxidizing environment. This is typically the case of coatings, whether metallic or ceramic, which form oxides such as alumina, chromia and silica. However, it is to be noted that inert barriers have been proposed in some cases, e.g. silica coatings on steels [7], iridium on C/C [ 1] and thin platinum layers on titanium alloys [8]. Their effective use would imply that a partial absence of coating, due for instance to a coating defect or spalling, is not too detrimental to the component life as the protection cannot then be restored in these systems. Moreover, interdiffusion effects between coating and substrates are supposed to be innocuous as far as the properties of the coated component are concerned.
2.1. Diffusion coatings 2.1.1. Simple diffusion coatings Aluminide coatings [9, 10] are the most widely used to protect nickel- and cobalt-based superalloys for gas turbine applications. They are generally formed by pack cementation: the components are immersed in a powder mixture in a semi-sealed retort. The retort is then placed inside a furnace and heated at high temperature (between 750 °C and 1150 °C) under a protective atmosphere. In the case of an aluminizing process, aluminium is transferred from the pack to the substrate surface where it reacts and an aluminide layer forms by diffusion with the substrate. Due to this mode of formation this technique is a particular ease of "reactive chemical vapour deposition (CVD)". Pack cementation is a simple process to operate, easily controlled and inexpensive to run. These advantages justify its wide use for protecting superalloys, but due to its mode of formation, it suffers from several limitations: the range of accessible coating compositions and structures remains rather narrow (only one element in practice can be transferred in the gas phase from the pack) and in some applications it would be desirable to enrich the substrate surface not only with aluminium, but also with chromium, yttrium, etc. Moreover, the composition and structure of the substrate alloy can seriously affect the performances of the coating. Other diffusion coatings have been developed: silicon carbide on carbon-based materials [ 11] and silicides on niobium alloys [ 12]. In the latter case the addition elements are brought to the surface as a slurry. 21.2. Modified diffusion coatings As stated earlier, the range of accessible compositions and structures with a pack cementation process is limited. However in several applications, it would be interesting to benefit from the advantages of reactive CVD processes (uniform thickness on complex geometries in particular) but to modify the coating composition in order to improve its corrosion resistance. Different routes have been proposed to modify aluminide coatings: by a pretreatment or a predeposit prior to the aluminizing process or by codeposition. A typical pretreatment consists in chromizing before aluminizing to obtain a chromium-enriched aluminide coating. Industrially operational chromium-modified aluminide coatings include C1A (SNECMA) and PWA 32.
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Codeposition of two (or more) elements during a single-step cementation treatment does not seem to be a widespread technique in industrial practice although it was envisaged a long time ago. Reports of combined chromium and silicon cementation of steel were published as early as 1945 [13]. Recent works based on thermodynamic calculations have demonstrated the feasibility of codepositing chromium and aluminium on nickel-base alloys [14, 16] or on iron-base alloys [15, 16], as well as aluminium and hafnium [16]. Difficulties seem to remain concerning the codeposition of aluminium and an active element such as yttrium, the solubility limit of this element being very low in nickel- or iron-based alloys. It is assumed that the available information concerning the transfer to industrial use (control of cement composition, etc.) is scarce. Incorporation of a third element by a predeposit has led to the development of platinummodified coatings. Most often platinum is electrodeposited as a layer 8/~m thick. After aluminizing, platinum is concentrated in the external region of the coating which is generally constituted of a mixture of PtA12 and fl-NiAl enriched in platinum. The addition of platinum improves
the resistance of simple fl-NiAl-based coatings, especially against type I hot corrosion. Recent work by Alprrine et al. [17] indicates that platinum may be replaced by the chemically similar palladium, which is less expensive, provided proper care is taken not to embrittle the coating during its processing
2.2. MCrAIY overlay coatings [18] MCrAIY alloys whose compositions and structures are entirely independent of the substrate can now be deposited as a result of the development of now well-established techniques such as electron beam physical vapour deposition (EB-PVD), and plasma spraying under a controlled atmosphere. As a consequence, a wide range of compositions is commercially available as shown in Table 1. 2.2.1. Structure and compositions MCrA1Y alloys are most generally multiphase materials. If the aluminium content is not too high, nickel- and cobalt-base alloys consist basically of a ductile ~ solid solution (face-centred cubic) containing a dispersion of //-NiAI or /%CoAl phase. The actual microstructure can be
TABLE 1 Examples of MCrAIY nominal compositions (wt.%) which are commercially available 119-211 Designation a
LCO ATD LCO ATD LCO ATD
Ni
29 5 5 2 7 14
Co
Fe
Bal. Bal. Bal. Bal. Bal. Bal.
Cr
AI
Y
18 18 19 23 23 30
8 11 10 12 13 5
0.5 0.3 0.5 0.3 0.6 0.5
17 22 25 25 38
6 10 12 5 11
0.5 1. 0.3 0.4 0.3
Amdry 961 Amdry 962 ATD 16 Amdry 963 ATD 1
Bal. Bal. Bal. Bal. Bal.
ATD 7 ATD 9 LN 21 Amdry 997 LN 34
Bal. Bal. Bal. Bal. Bal.
22 20 22 23 0.5
18 42 21 20 20
12 5 7.5 8.5 11
0.3 0.3 0.5 0.6 0.5
LCO 22 Amdry 995 LCO 37 Amdry 996
32 32 23 10
Bal. Bal. Bal. Bal.
21 21 30 25
7.5 8 3 7
0.5 0.5 0.5 0.6
Amdry 970 ATD 8
Bal. Bal.
24 24
8 11
Other
4 Ta 0.5 Mo
5 Ta
0.4 0.6
aATD = coatings produced by EB-PVD, Airco Temescal; LN and LCO = coatings produced by argon-shield plasma spraying, Union Carbide; Amdry = powders available for LPPS, Alloy Metals.
16
more complex as shown for example by Frances et al. [22] who have identified ~', o, M'sY and Y203 phases in addition to fl and ~ in NiCoCrAIYTa alloys. It is to be borne in mind that the microstructure of MCrAIY alloys may depend on previous heat treatments and cooling rates. If the aluminium content is sufficiently high, the fl phase percolates and contains a dispersion of ~-Ni particles. Such a structure is certainly less favourable in terms of ductility. It is to be noted that in the Ni-Cr-A1 ternary system, a phase transformation can occur at about 1000 °C whereby ~ + fl, stable at high temperature, transforms into y'-Ni3Al+ ~-Cr as represented schematically on Fig. 1 based on the phase diagrams established by Taylor and Floyd [24]. This reaction is accompanied by a significant volume variation (Fig. l(b)) which may be deleterious for the mechanical integrity of a coatingsubstrate system. It is therefore recommended that an alloy composition outside this field is selected, limiting the aluminium content, or that cobalt is added to destabilize ~' formation. The oxidation behaviour of MCrAIY-coating alloys depends on various factors (for a review see, for example, ref. 25), e.g. alloy composition, temperature, oxygen partial pressure, time and thermal cycling. A typical oxidation sequence for an MCrA1Y alloy starts with the formation of transient oxides (NiO, spinel, etc.) growing rapidly until a continuous stable oxide layer (chromia or alumina) has formed underneath. The diffusion of oxygen and metallic species being very slow in chromia and even slower in alumina, this layer ensures the oxidation resistance of the alloy. The presence of chromium in these alloys reduces the aluminium level necessary to form a protective alumina scale. The presence of as little as 5 % - 1 0 % Cr reduces
N~62Cr21AII7
(a)
..... .
(b)
Fig. 1. (a) Partial phase diagram in the N i - C r - A 1 system and (b) thermal expansion data showing the phase transformation at around 1000 °C [23].
the amount of aluminium required from 40 to about 10 at.%. Alumina scales formed on MCrA1 alloys are not adherent in thermal cycling conditions, presumably as a result of growth and thermal stresses. In order to improve the adherence of the alumina scale, so-called active elements such as yttrium, rare-earths, etc. are incorporated in these alloys in small amounts (less than 1 wt.%). Several explanations have been put forward to explain this effect but a controversy still exists relative to the precise mechanism(s) involved. Other elements (e.g. hafnium and zirconium) have been claimed to have the same effect. A recent comparative study [23] showed however that, of hafnium, zirconium and yttrium, the latter is by far the most effective addition element for improving the adherence of alumina scales formed at 1100 °C. It is to be noted that this effect is observed whether yttrium is initially present in the oxide particles or metallic phases within the coating. It seems in fact that a uniform distribution of these elements is essential. Recently, Gupta and Duvall [26] recommended the addition of both yttrium and hafnium. A major degradation (termed "hot corrosion") is the accelerated oxidation induced by condensed salts formed by pollutants present in combustion gases. Major contaminants in gas turbines are alkaline sulphates, in particular Na2SO4. The degradation modes depend on various factors: temperature range, nature of contaminants, gas composition, alloy composition [5]. Several mechanisms have been proposed (for a review, see ref. 27). At relatively high temperature (type I hot corrosion, T > 800 °C), the degradation of the scale can take place either by basic fluxing, which affects alumina-forming alloys containing less than about 15% Cr or by acidic fluxing induced by oxides such as MOO3, WO3, V205 formed by the oxidation of refractory elements. At moderate temperatures (650 °C < T < 800 °C) another form of hot corrosion (type II [28]), which affects alloys having a low chromium content, has been reported. It involves acidic fluxing of oxide scales by SO3 dissolved in molten sulphates in the presence of air. The presence of chromium is beneficial against all types of hot corrosion. It is thought that Cr203 acts as a buffer against basic fluxing mechanisms. Moreover, Cr203 is formed more rapidly than AI203. About 2 5 % - 4 0 % Cr is required for type II hot corrosion whereas 15%-20% Cr is usually
17 recommended for type I [28]. It is to be noted that too high a chromium content may have deleterious effects on the diffusional stability of the system. Silicon is particularly good in the case of type II hot corrosion (acidic fluxing) but above about 800 °C, it tends to diffuse rapidly into the substrate and may form low-melting phases or a brittle intermetallic compound. Other addition elements may play a beneficial role; tantalum, for example, although not particularly good for oxidation resistance [29], prevents the outward diffusion of potentially deleterious elements such as titanium by tying them up, with carbon, in stable MC carbides. The above considerations justify the wide use of standard NiCoCrAIY-base compositions with about 20% Co, 20% Cr, 8 % - 1 2 % AI and < 1% Y(wt.%). 2.2.2. Alternative deposition techniques for metallic coatings
Despite the flexibility they permit, the techniques commercially available for depositing MCrAIY coatings, EB-PVD and plasma spraying under inert atmosphere, remain line-of-sight processes and this, in addition to their relatively high cost, can be a real drawback for coating components having complex shapes. In view of some of the disadvantages of production techniques, alternative processes have been developed, among which electrolytic codeposition and electrophoresis seem the most promising at the moment. 2 . 2 . l 1. Electrolytic codeposition. MCrA1Y coatings have been obtained by co-electrodeposition of a dispersion of fine CrA1Y powder particles and a cobalt or nickel matrix [30]. This operation can be carried out in a barrel plating unit (Fig. 2). The barrel containing the specimens and the powder particles (10 #m diameter) is immersed in a cobalt and/or nickel conventional
Drive
\ ~ 9ear
Sol~
,n.i
level
(~
-i
Su
r
Drive shaft
Flexible
@ rotating Contact holder
I
~
(tornotor}
bath and rotated. The coated components are subsequently heat treated for alloying and diffusion bonding the coating to the substrate. This process presents the obvious economical advantages of low capital cost (claimed to be an order of magnitude less than existing commercial systems; unit processing cost about half those of competing systems [30]). Results of reliability tests do not seem to have been published as this technique has not yet been scaled up.
2.2.2.2. Electrophoresis. This process is based on the migration of fine powder particles suspended in a polar solution under a d.c. voltage. In the process developed at SNECMA [31,32], MCrA1Y alloy powders with size less than about 40 #m are deposited by electrophoresis on the components to be coated with deposition rates up to 100 #m min -~. Figure 3(a) shows the deposit with controlled porosity obtained at this stage. In order to densify them and adjust the coating composition, the components are then aluminized (vapour-phase process). The appearance of an MCrAIY coating obtained is illustrated in Fig. 3(b). With adequate electrode geometry, components with complex shapes can be uniformly coated. The facilities required are relatively simple, so the investment, as well as the running costs, is much less than for processes such as EB-PVD and plasma spraying. As shown in Fig. 3(c), these coatings confer a hot corrosion resistance comparable with that of low-pressure plasma-sprayed coatings. 2.2.2.3. Cladding. Cladding MCrA1Y sheets has been expected to produce protective coatings [33] but early experiments do not seem to have been an overwhelming success. This may have been caused by the very limited aluminium content of available sheet alloys (conventional forming processes require a low aluminium content) and possibly an inadequate yttrium distribution. It might be interesting to reconsider this technique since MCrAIY ribbons (50-100 pm thick) having more oxidation-resistant compositions can now be obtained by melt-spinning [34].
gear
2.3. Thermal barrier coatings (b
ath~e
Bearin9 /
Fig. 2. Schematic diagram of a co-electrodeposition plating unit [30]. (a) Longitudinal section; (b) transverse section.
Thermal barrier coatings (TBCs) generally consist of an oxide layer (thickness ranging from about 300 #m in turbine applications up to about 2 mm in diesel engines) plasma sprayed on top of an MCrA1Y bond layer deposited on the substrate.
18
(a) Electrophoresis + aluminizing (C3A)
1000
MCrAIY
(L.P.P.S,
1000
Cr + AI
(C1A)
Aluminide (A.P.V.S.) Uncoated
I zoo -715o
-]so Life time (h.)
Ba-
Burner rig (870°C)
(b)
(c)
Fig. 3. (a) Cross-section of an MCrA1Y coating after electrophoresis and (b) subsequent consolidation by vapour phase aluminizing, (c) hot corrosion resistance in burner rig tests, with Na2SO4 deposits, at 900°C of various coatings on IN 100 (c) [18].
The benefits of thermal barrier coatings are multiple: higher inlet gas temperatures, lower cooling flow within the metallic component, reduction of temperature transients on the metal surface. All these advantages can result in improved efficiency, extended lifetimes and simplified designs. Such ceramic coatings have been plasma sprayed on sheet metal combustor components for more than 15 years. Only recently [35] have they been used in the turbine section of commercial gas turbine engines. Thermal barrier coatings over the entire airfoil of PW4000 turbine vanes were certified in the U.S.A. for production use in 1986. This progress could be achieved by a proper control of porosity, microcrack distribution, and residual stresses in the coating. With nickel-base superalloys having reached their optimum properties, one of the present challenges t o increase gas turbine performance significantly remains the application of these coatings on highly stressed components such as turbine blades.
completely stabilized with proper additions of Y203, MgO or CaO, since ZrO2 offers a good compromise between low thermal conductivity and high thermal expansion coefficient. Yttria partially stabilized zirconia containing between 6 and 8 wt.% Y203 has been found empirically [36]
2.3. I. Zr02- (6-8)wt. % }'2 03 Figure 4 illustrates the typical aspect of the cross-section of a thermal barrier coating. Most often, the oxide is zirconia-based, partially or
Fig. 4. Typical appearance of the cross-section of a TBC showing partially stabilized zirconia (low thermal conductivity and high compressive strain tolerance), MCrA1Y bond coat (strain isolation and oxidation resistance) and superalloy substrate.
PARTIALLY STABILIZED ./ZIRCONIA
MCrAIY BOND COAT
SUPERALLOY SUBSTRATE
19 to be most resistant in thermal cycling conditions. This composition corresponds to the formation of a metastable tetragonal t' phase when the cooling rate is sufficiently high, as in plasma spraying, which is at the moment the deposition process most often used for obtaining these coatings. With proper parameter adjustment, plasma spraying gives a microcracked structure particularly strain tolerant in compression [37]. This is important as the main degradation mode (spallation) of TBCs results from fatigue cracking damage in the ceramic layer near the ceramic-bond coat interface induced by thermal strains. It should be emphasized though that the precise role of the t' phase in terms of crack propagation rate is not clearly understood and basic studies are needed to relate the fine microstructure and phase transformations of TBCs with their mechanical properties [38, 39]. Recent engine tests by Sheffler and Gupta [40] and Toriz et al. [41] have proved that TBCs obtained by EB-PVD may perform even better than plasma-sprayed coatings. This might be due to the well-adapted anisotropic microcrack network in the former case. 2.3.2. Role of the bond coat Between the oxide topcoat and substrate, an MCrA1Y bond coat confer an adequate environmental to adapt the oxide layer and the chanically.
the superalloy is required to resistance and substrate me-
2.3.2.1. Mechanical adaptation. This mechanical adaptation is more complex than a simple matching of thermal expansion coefficients. Actually the expansion coefficients of MCrAIY alloys are generally higher than for superalloys. The calculation of residual stresses in a purely elastic situation (with experimental elastic moduli) gives values well above the measured stresses. In fact, strain accommodation occurs to a large extent by plastic deformation of the bond coat as MCrAIY alloys have a very low flow stress at temperatures above 600 °C. 2.3.2.2. Environmental resistance. Thermal barrier coatings are dynamic systems in the sense that zirconia top coat, MCrA1Y bond coat and substrate superalloy interact at high temperature. Figure 5 shows, for example, the thin alumina layer formed at the zirconia-MCrA1Y interface after a 100 h heat treatment at 1100 °C. This layer has
Yttrim partial#y stabilized zirconia
Alumina
NiCrAIY
.
.
.
.
Fig. 5. Aluminalayer formedat the zirconia-MCrAlYinterface after 100h at 1100°C. formed by interaction of aluminium which has diffused from the MCrAIY alloys with oxygen present in the zirconia material. It is interesting to note that the adhesion between these layers (zirconia-alumina-MCrA1Y) is sound and that, in practice, spallation rarely occurs at the interfaces. Instead, in thermal cycling tests, failure of the TBC happens as shown on Fig. 6(a) by propagation of a main crack inside the zirconia top coat near the alumina-zirconia interface. Propagation results from compressive stresses arising in the top coat on cooling because of a thermal expansion mismatch. A more extensive oxidation of the bond coat can lead to the destruction of the barrier as shown in Fig. 6(b) as the transformation of the alloy into oxide is accompanied by a volume expansion which generates additional stresses in the zirconia layer. The evolution of the structure and composition of the MCrA1Y bond coat are likely to affect its mechanical properties and, as a consequence, to alter its ability to deform rapidly during cooling. This in turn will raise the compressive stress level in the ceramic layer, accelerating the degradation of the coating. Another degradation mode occurs by hot corrosion with liquid sulphate contaminants as illustrated in Fig. 7 [42]; when tested in cyclic oxidation at 950 °C, the weight variations are regular and no degradation occurs before 1000 cycles (test duration); however when tested in the same conditions with additional sodium sulphate contamination, a rapid degradation can be observed after only 50 cycles. Several mechanisms have been proposed to explain the accelerated degradation of the TBC in hot corrosion conditions: destabilization of zirconia, sulphidation of the bond coat, solidification of sodium sulphate in the cracks [42]. In the case of Fig. 7 these mechanisms did not seem to be operative and the
20 Yttria partially stabilized zirconia
Fig. 6. Typical degradation of a TBC: (a) propagation of a main crack parallel to the ceramic-metal interface and (b) internal oxidation of bond coat. Weight gain (mg/cm2) ~SALT
t
~/~/~ SOLIDIFICATION \
I NoContamination
01 -1 -2
: : :
\
' !
~ SALT ~ EVAPORATION~:
%.
-3
\ \
r BOND // COATING / OXIDATION
,3
o
200 400 6oo Boo Number of one hour cycles
1000
Fig. 7. Weight variations of a TBC tested at 950 °C in cyclic oxidation and in hot corrosion [42].
~MOLTENFD LIM AMAGESALI
ZIRCONIA~ TOUGHNESS ,~ REDUCTIONJ"
TEMPERATURE
Fig. 8. Schematic representation of TBC life as a function of temperature [43, 44].
explanation put forward involves a crack propagation ass~ted by the sulphate vapour.
2.3.3. Modelling A significant effort at modelling the behaviour and life of thermal barrier coatings is going on in the U.S.A., as summarized by Miller [43]. The models developed result from a predominantly phenomenological approach schematically illustrated in Fig. 8 and involve several empirical parameters. It must be emphasized that, up to now, some areas are still generally poorly described: creep and inelasticity of ceramic and bond coat, geometry effect (roughness of the metal-ceramic interface), oxidation-induced damages, sintering of oxide at high temperatures, etc. 2.3.4. Concluding remarks Thermal barrier coatings should provide the next significant increase in engine performance; Shetfler and Gupta [40] report that a zirconia
layer 250/zm thick can reduce the metal temperature by as much as 170 °C. This may be achieved through the development of more reliable coatings and of thermomechanical modelling. Moreover new ceramic compositions are required for improved coatings resistant to corrosive environments with contaminated fuels and molten salt deposits [45] and, in the long run, for coatings having a higher temperature capability than yttria partially stabilized zircouia. Finally new processes for depositing ceramic coatings will have to be developed to protect components having complex geometries.
3. Compatibility considerations Numerous considerations must be taken into account when selecting or designing a coating for a particular application; they are related to the
21
environment, the coating process, the service conditions and the substrate. Due to the drive for ever-increasing temperatures to improve engine performance and the development of advanced superalloys, considerations of coating-substrate compatibility, whether chemical or mechanical, become more and more important.
3.1. Chemical compatibility Coating and superalloy substrates generally have widely different compositions, each of them selected for different purposes: protection against oxidizing environment on the one hand and hightemperature mechanical properties on the other. The chemical potential gradients of the different elements (e.g. nickel, aluminium, chromium) across the coating-substrate interface will cause interdiffusion phenomena to occur and these effects will be all the more important when .the temperature is elevated. As a consequence, elements from the substrate may diffuse through the coating, reach the external surface and perturb the formation of the protective oxide scale. This is illustrated in Fig. 9 which represents the weight variations in cyclic oxidation at l l00°C of different superalloys (IN100, IN738 and CMSX2) coated with the same plasma-sprayed CoCrAIY alloy. The graph clearly shows that after only 250 cycles the coating deposited on IN100 and IN738 suffers extensive damage, which corresponds to internal oxidation and subsequent scale spallation, as revealed by metallographic examination, whereas on CMSX2 the coating lifetime is more than 1000 cycles. Analysis of oxidized samples suggested that, in the first two cases, the formation of the alumina scale is perturbed by diffusion of titanium from the substrates toward the external surface.
5
Am mg cn~ 2
3 2 l
Number Of cycles exposure time at HO0°C
SX2
Fig. 9. Cyclic oxidation at l l 0 0 ° C of a CoCrAIY coating low-pressure plasma-sprayed on different superalloys [46].
Conversely, interdiffusion may also affect the structure of the subcoating zone in the substrate and therefore alter or modify its mechanical properties. Several transformations can take place: (a) transformation of the initial y - y ' microstructure into y - 8 in the case of CoCrAIY/Nibase superalloys [46, 47], (b) alteration of the y' distribution, (c) modification of the distribution of grainboundary strengthening elements (carbon, etc.) [48], (d) destabilization of the reinforcing carbide fibres in directionally solidified eutectics [49], (e) formation of brittle phases which may be crack initiation sites or easy paths for crack propagation. These effects, which can be important in the case of components with thin sections, can be minimized by carefully selecting a coating composition compatible with that of the substrate. This is possible in the case of MCrAIY coatings because of the flexibility allowed by the deposition processes. Some authors [50] have suggested the interposition of a "diffusion barrier" (generally a stable ceramic material in which metallic diffusion is slow) between coating and substrate to prevent these effects. It is felt, however, that such an intermediate layer would present several drawbacks: poor thermomechanical behaviour due to thermal expansion mismatch and a preferential path for oxidation penetration once the diffusion barrier material has been reached by oxygen through a coating defect initially present or induced by local degradation in service.
3.2. Mechanical compatibility A protective coating may alter the mechanical properties of a superalloy substrate on which it is applied as a result of factors given above. It must be stressed, however, that in the long run, a coating, if adequately protective, will isolate the superalloy from the environment (oxidizing combustion gases, molten salt deposits) and will limit or suppress the corrosion-induced mechanical degradation. In a study on thin-section Ren6 80 specimens precorroded for 500 h in an engine atmosphere, Kaufman [51] has shown that the presence of a coating, whether aluminide or MCrAIY, improves the creep lifetime (at 980 and 1080 °C) by factors ranging between 2 and 6, as well as the fatigue resistance at room temperature (220-280 MPa).
22
It is generally accepted that creep properties of superalloys are not affected by a coating provided the standard ~heat treatments are respected and little interdiffusion occurs in service or during the application of the coating (Fig. 10); a possible favourable influence can occur in the long run because of the suppression of environmental degradation. The mere presence of a protective coating can significantly alter the fatigue resistance of superalloy substrates. This has been clearly shown, for example, by Veys and M6vrel [52] and by Wood and Restall [53, 54] in the case of single-crystal superalloys (Fig. 10). The reduction in fatigue lifetime is related:to early cracking in the coating with subsequent propagation inside the substrate. The fact that early cracking occurs at temperatures around 800 °C or 900 °C may be an aggravating factor for hot corrosion degradation: in this temperature range, sodium sulphate may condense o n blade surfaces and penetrate through these.cracks, abolishing the beneficial effect of the protective coating. It is important to note that the structure and composition of a protective coating change during operation a t high temperatures as a result of aluminium consumption by oxidation; the initial oR (MPa) I N -x"~- iCMSX'2unc°ated[11] •~ : CMSX-2/alurninide : CMSX-2/N CoCrA TaY 500 ! 400.
e'~e
300
200
\
\
~ ' ex
150 1'20"T(OK).10?[20 + log tr,] =25 2'9 30 :Jl Omax (MPa) o \ X~X 800
0 A
UncoatedCMSX-2 [11]
o
CMSX-2 + aluminide
¢
CMSX:2+ NiCoCrA YTa T =" 870°C = 70 HZ
" oN ~ ~ El
[3
70~ 0
.6.0p 0.1 "
1
"1'0
"Nr (cycles) . 106 100 =
Fig'. 10~ Creep and high cycle fatigue properties o f coated CMSX-2 [52].
+/~ structure of an MCrAIY coating therefore progressively transforms into ~ ( + ~ ' ) and in the end to y solid solution of nickel; this is accompanied by grain growth. No study however seems to have been published on these effects, the durations of tests performed being too short for these phenomena to be significant. Recent efforts [55-58] have concentrated on thermomechanical fatigue tests which combine mechanical strain-controlled cycling and temperature cycling; these tests offer a better simulation of the complex strain-temperature-time cycles experienced by gas engine components in service. Moreover, thermomechanical fatigue may now b e a major limiting factor in turbine vanes and blades [37], in particular for highly stressed thinwalled components. Swanson et al. [57, 58] are developing life prediction models with the emphasis on applicability to thermomechanical fatigue conditions. These models are based on the development of individual constitutive models of coatings (PWA286: NiCoCrAIYSiHf and PWA273 low activity aluminide) and PWA1480 single-crystal substrate, for which commercial production experience already exists. Results of out-of-phase tests carried out between 427 and 1038 °C indicate that at strain ranges greater than 0.5%, the overlay-coated specimens have longer lives than do those coated with the aluminide, the reverse being true at lower strain range. It is interesting to note that similar results: had been obtained by Bain [56] on a single,crystal superalloy coated with a CoCrA1Y alloy or an aluminide and tested in slightly different conditions. As emphasized by Wood [59], little work has been done to determine in depth which coating properties are of importance in controlling the mechanical behaviour of coating-substrate systems. The studies in this field have dealt mainly with the identification of degradation modes and the experimental evaluation of the incidence of coatings on a few mechanical properties of ascoated substrates. It is important to point out that coatings are found as crack initiation sites in a variety of test conditions, the cracks thus formed propagating subsequently into the substrate material and leading eventually to failure. Neither the evolution of the coating composition in service nor the presence of corrosive products are taken into account. Some attempts at modelling the coating-superalloy system have been recently published [58]. However, the lack of data relative
23 to the intrinsic mechanical properties of the coating, mainly in the case of the widespread aluminides (simple or modified), seems t o have delayed the development of an adequate description of these systems. As a consequence, the possibility of tailoring the Coating composition and/or structure to improve the mechanical properties of the coating-substrate composite does not seem to have been fully exploited.
Acknowledgments Thanks are due to m a n y people for valuable discussions and suggestions, in particular Serge Alp6rine and Pierre Jesse on palladium-modified aluminide coatings, Patrick Choquet on the oxidation of MCrA1 + X coatings, J,-M. Veys on the mechanical properties of coated systems, Ren6 Morbioli ( S N E C M A ) on electrophoretic deposits and Robert Miller (NASA) on thermal barrier coatings.
4. Conclusion Looking back on the historical evolution of high-temperature protective coatings it is interesting to note that the systems with the best inservice performance at the m o m e n t have been developed according to a largely empirical approach. Further developments can be expected along this line, for example with the introduction of alternative coating processes for depositing MCrA1Y alloys (electrophoresis, co-electrodeposition) and with the replacement of platinum by palladium in modified aluminide coatings. It has to be realized, however, that the development of more advanced processes (codeposition in single step cementation for instance) requires unavoidable preliminary theoretical treatments. More indepth investigations are needed in order to understand the degradation mechanisms better and therefore to be in a position to exploit the full potential of available coating compositions. In this respect more data and modelling are necessary concerning the intrinsic mechanical properties of coatings and also the thermomechanical behaviour of coated systems. Application of thermal barrier coatings on highly stressed components represents the next challenge for significantly improving gas turbine engine performances. This implies an active research effort in the direction of more reliable coating processes, and, depending on the applications envisaged, of new systems with higher temperature capabilities and improved hot corrosion resistance. Finally an increasing activity is noted concerning the development of high-temperature materials: titanium alloys having creep strength at temperatures above those for which oxidation resistance is adequate, intermetallic compounds, etc. Formulating new coating compositions and processes is expected to be necessary to ensure the protection of these materials.
References 1 J. R. Strife and J. E. Sheeban, Ceram. Bull., 67 (~988) 369-374. 2 G. W. Goward, Mater. Sci. Technol., 2 (3) (1986) 194200. 3 J. Stringer, in R. Kossowsky and S. C. Singhal (eds.), Surface Engineering, NATO AS1 Ser. 85, pp. 561-587. 4 R. A. Miller, Surf. Coat. Technol., 30 (1987) 1-11. 5 C. Duret-Thual, R. Morbioli and P. Steinmetz, A Guide to the Control of High Temperature Corrosion and Protection of Gas Turbine Materials, Commission of the European Communities, EUR 10682, 1986. 6 High Temperature Oxidation Resistant Coatings, National
Academy of Sciences and Engineering, 1970, 7 M. J. Bennett, J. Vac. Sci. TechnoL B, 2 (1984) 800-805. 8 S. Fujishiro and D. Eylon, Thin Solid Films, 54 (1978) 309-315. 9 R. Mtvrel, C. Duret and R. Pichoir, Mater Sci. Technol., 2 (1986) 201-206. 10 R. Mtvrel and R. Pichoir, Mater. Sci. Eng., 88 (1987) 1-9. 11 M.-P Bacos, personal communication, 1989. 12 S. Priceman and L. Sama, Electrochem. Technol., 6 (9-10) (1986) 315-326. 13 Anon. Iron Age (October, 1945) 58-61. This article gives an abstract and translation of a report published by A. M. Borzdika, StaL 7-8. 14 G. H. Marijnissen, in S. C. Singhal (ed.), High Temperature Protective Coatings, The Metallurgical Society of AIME, Warrendale, PA, 1983, pp. 27-35. 15 R. A. Rapp, D. Wang and T. Weisert, Prec. TMS-AIME Symp. Metallurgical Coatings, Orlando, 1986, The Metallurgical Society of AIME, Warrendale, PA. 16 B. Nciri, Thdse, Orltans, 1983. G. Leprince, S. Alptrine, L. Vandenbulcke and A. Walder, Mater. Sci. Eng., A120/121 (2) (1989) 419. 17 S. Alptrine, P. Steinmetz, P. Jesse and A. Costantini, Mater. Sci. Eng., A120/121 (2) (1989) 367. 18 R. M6vrel and R. Morbioli, in Prec. Ist Int. Congr. on High-Tech-Materials and Finishing, Berlin, March, 1989,
pp. 205-216. 19 D. H. Boone, Mater. Sci. Technol., 2 (3) (1986) 220224. 20 B. J. Gill and R. C. Tucker, Mater. Sci. Technol., 2 (3) (1986) 207-213. 21 AMDRY plasma and flame spray powders, Tech. Note. 22 M. Frances, M. Vilasi, M. Mansour-Gabr, J. Steinmetz and P. Steinmetz, Mater. Sci. Eng., 88 (1987) 89-96. 23 P. Choquet, Thdse Doctorat ds Sciences, Orsay, 1987. 24 A. Taylor and R. W. Floyd, J. Inst. Met., 81 (1952-3) 451-464. 25 F. H. Stott, Rep. Prog. Phys., 50(1987) 861-913.
24 26 D. K. Gupta and D. S. Duvall, in M. G-ell et al. (eds.), Superalloy 1984, The Metallurgical Society of AIME, Warrendale, PA, 1984, pp. 713-720. 27 R. A. Rapp, Mater. ScL Eng., 87(1987) 319-327. 28 G. W. Goward, ASME Pap., 85-GT-60(1985) 1-5. 29 M. Vilasi, Thdse, Nancy, 1988. 30 F. J. Honey, E. C. Kedward and V. Wride, J. Vac. Sci. Technol. A, 4 (6) (1986) 2593-2597. 31 Aboulicam, Gauj6, Grammagnac and Morbioli, Demande de brevet d'invention 2 529 911, 1982. 32 R. Morbioli, Mater. Sci. Eng., A120/121 (2) (1989) 373. 33 Final Rep. EPRI AP-3267 (1983). 34 F. Duflos, C. Duret, V. Gossart, M. Kumar and A. Walder, in High Temperature Alloys for Gas Turbines and Other Applications 1986, Part II, Reidel, Dordrecht, 1986, pp. 989-998. 35 J. W. Fairbanks and R. J. Hecht, Mater. Sci. Eng., 88 (1987) 321-330. 36 S. Stecura, NASA Tech. Memo. TM-86905(1985). 37 "I~. A. Cruse, S. E. Stewart and M. Ortiz, Trans. ASME J. Eng. Gas Turbines Power, 110 (1988) 610-616. 38 R. McPherson, Surf. Coat. Technol., in press. 39 L. Lelait, S. All~rine, C. Diot and R. M6vrel, Mater. Sci. Eng., A120/121 (2) (1989) 475. 40 K. D. Sheffler and D. K. Gupta, Trans. ASME J. Eng. Gas Turbines Power, 110 (1988) 605-609. 41 F.C. Toriz, A. B. Thakker and S. K. Gupta, to be published. 42 S. Alp6rine, AGARD SMP Meeting, Ottawa, 1989, to be published. 43 R. A. Miller, NASA Tech. Memo. TM-100283 (1988). 44 T. E. Strangman, J. Neumann and A. Liu, NASA (1987) Contract. Rep. CR-179648. 45 R. L. Jones, Paper presented at ICMC 1989, San Diego, CA, 1989, to be published. 46 R. M6vrel and C. Duret, in Coatings for Heat Engines,
NATO Advanced Workshop, Acquafredda di Maratea, 1984, pp. 595-612.
47 P. Mazars, D. Manesse and C. Lopvet, Interdiffusion of MCrA1Y coatings with substrates, in High Temperature Alloys for Gas Turbines and Other Applications 1986, Reidei, Dordrecht, 1986, pp. 1183-1192. 48 A. Strang, E. Lang and R. Pichoir, AGARD SMP, CP356, 11/1-11/35, 1983. 49 J.-M. Veys and R. M6vrel, Mater. Sci. Eng., 88 (1988) 253-260. 50 J. P. Coad, D. S. Rickerby and B. C. Oberlander, Mater. Sci. Eng., 74 (1985) 93-103. 51 M. Kaufman, Examination of the influence of coatings of thin superaUoy sections, NASA Contract. Rep. CR-1215 (1972). 52 J.-M. Veys and R. M6vrel, ASM Technical Conference, Paris, September, 1987, American Society for Metals, Metals Park, OH. 53 M. I, Wood and J. E. Restall, in High Temperature Alloys for Gas Turbines and Other Applications, Liege, 1986, pp. 1215-1226. 54 M. I. Wood, Proc. 1st ASM Europe Technical Conference on Advanced Materials and Processing Techniquesfor Structural Applications, Paris, September 1987, ONERA, 1987, pp. 179-188. 55 T. E. Strangman, PhD Thesis, University of Connecticut, 1978. 56 K. R. Bain, AIAA-85-1366, 1985, pp. 1-6. 57 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, NASA Contract. Rep. CR 174952 (1986). 58 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, NASA Contract. Rep. CR 179594 (1987). 59 M. I. Wood, Mater. Sci. Eng., A120/121 (2) (1989) 633.
13
State of the Art on High-temperature Corrosion-resistant Coatings R. M]~VREL ONERA, Materials Science Department, BP 72, 92322Chtitillon (France) (Received July 10, 1989)
Abstract
A high-temperature protective coating must meet several criteria: provide adequate environmental resistance, be chemically and mechanically compatible with the substrate, be applicable. Comprehensive reviews on high-temperature coatings have appeared regularly since the early 1970s. Our purpose is not to recapitulate the material covered therein but rather to focus on recent trends, and point out some research perspectives. Historically the development of high-temperature protective coatings has been linked with the evolution of demanding applications such as superalloy components in gas turbine engines; the searches for better performance (higher inlet temperatures, longer lifetimes, etc.) and for cost-saving solutions (use of contaminated low-grade fuels) have been the main incentives for developing the different coatings now available in production: simple and "modified" diffusion coatings, overlay coatings, thermal barriers. In recent years research and development activity has been concentrating on the following points. (a) Degradation mechanisms in high-temperature corrosion of metallic coatings; basic studies on the growth mechanisms of oxide scales, for exampie, are still required, in particular to understand the role of addition elements such as platinum and palladium. (b) Alternative techniques for depositing MCrAIY coatings; electrolytic codeposition and electrophoresis, for example, have been developed at the laboratory stage and these permit the deposition of MCrAIY coatings with claimed economic and technical advantages over processes already in production. (c) Thermal barrier coatings; ceramic coatings have been applied to sheet metal combustor components for about 15 years; only recently have they been used in the turbine section. Two challenges remain though to exploit these coatings on turbine blades: improve their reliability and, in the case of 0921-5093/89/$3.50
stationary gas turbines, their hot corrosion resistance. Both structural and mechanical approaches are required to determine, in particular, the role of microstructure, microcracking, porosity, residual stresses, oxidation of the bond layer in the degradation mechanisms of these coatings. (d) Mechanical properties of coated systems; the intrinsic mechanical properties of coating materials are still poorly described and the lack of information hinders the adequate modelling of the behaviour of coating-substrate composite systems. In parallel, an increasing activity is noted concerning the design and development of high-temperature coatings for protecting materials other than superalloys, for instance ceramic composites and titanium-based alloys.
1. Introduction
Materials for high-temperature applications are generally selected for specific properties such as strength, creep, mechanical and/or thermal fatigue. During their use at high temperatures these properties may degrade as a result of interactions with corrosive environments. One solution to avoid this alteration is to protect the component with a coating. Most high-temperature coatings rely on the formation of a protective oxide scale by interaction with the environment. The role of this scale is to isolate the coating from the aggressive environment and thereby limit the degradations due to high-temperature corrosion. Obviously this scale must fulfil several conditions: be stable, slow-growing, dense, adherent (often in cyclic conditions). In practical applications only three oxides correspond to these criteria: alumina (A1203) , chromia (Cr203) and silica (SiO2). The fact that chromia can form volatile suboxides at temperatures above about 900 °C generally restricts its use to somewhat lower temperatures, especially with high-velocity gases. Silica scales © ElsevierSequoia/Printed in The Netherlands
14 confer protection on ceramic coatings (SIC, Si3N 4 for example) deposited on carbon-based materials up to about 1800°C [1]. However, the use of silica-forming metallic alloys as coatings has been limited. This is due to the fact that in order to form a silica scale, a metallic coating must contain a relatively large amount of silicon. But silicon, which diffuses rapidly in most alloys, may form low-melting phases and/or brittle silicides by interaction with substrates. As in the case of chromia-forming coating alloys, the use of silicaforming metallic coatings must be restricted to rather low temperatures. These considerations explain why most high-temperature coatings (simple or modified aluminides, MCrA1Y overlays) rely on the formation of an alumina scale to ensure the protection of an underlying component. Comprehensive reviews on high-temperature protective coatings have regularly appeared since the early 1970s [1-6]. Our purpose is not to recapitulate the material covered therein, but rather, on the basis of published material and some personal communications, to focus on recent trends and to attempt to point out some research perspectives in the first part of this paper. The second part is devoted to substrate-coating compatibility considerations. 2. Main types of coating As mentioned above, most coatings rely on the isolation by a "barrier" of the part to be protected from the environment. The role of this barrier is to minimize the diffusion of gaseous or liquid species towards the component and conversely to prevent elemental diffusion from the alloy towards the external surface where they could react. Most generally the barrier is dynamic in the sense that it forms by interaction of the coating with the oxidizing environment. This is typically the case of coatings, whether metallic or ceramic, which form oxides such as alumina, chromia and silica. However, it is to be noted that inert barriers have been proposed in some cases, e.g. silica coatings on steels [7], iridium on C/C [ 1] and thin platinum layers on titanium alloys [8]. Their effective use would imply that a partial absence of coating, due for instance to a coating defect or spalling, is not too detrimental to the component life as the protection cannot then be restored in these systems. Moreover, interdiffusion effects between coating and substrates are supposed to be innocuous as far as the properties of the coated component are concerned.
2.1. Diffusion coatings 2.1.1. Simple diffusion coatings Aluminide coatings [9, 10] are the most widely used to protect nickel- and cobalt-based superalloys for gas turbine applications. They are generally formed by pack cementation: the components are immersed in a powder mixture in a semi-sealed retort. The retort is then placed inside a furnace and heated at high temperature (between 750 °C and 1150 °C) under a protective atmosphere. In the case of an aluminizing process, aluminium is transferred from the pack to the substrate surface where it reacts and an aluminide layer forms by diffusion with the substrate. Due to this mode of formation this technique is a particular ease of "reactive chemical vapour deposition (CVD)". Pack cementation is a simple process to operate, easily controlled and inexpensive to run. These advantages justify its wide use for protecting superalloys, but due to its mode of formation, it suffers from several limitations: the range of accessible coating compositions and structures remains rather narrow (only one element in practice can be transferred in the gas phase from the pack) and in some applications it would be desirable to enrich the substrate surface not only with aluminium, but also with chromium, yttrium, etc. Moreover, the composition and structure of the substrate alloy can seriously affect the performances of the coating. Other diffusion coatings have been developed: silicon carbide on carbon-based materials [ 11] and silicides on niobium alloys [ 12]. In the latter case the addition elements are brought to the surface as a slurry. 21.2. Modified diffusion coatings As stated earlier, the range of accessible compositions and structures with a pack cementation process is limited. However in several applications, it would be interesting to benefit from the advantages of reactive CVD processes (uniform thickness on complex geometries in particular) but to modify the coating composition in order to improve its corrosion resistance. Different routes have been proposed to modify aluminide coatings: by a pretreatment or a predeposit prior to the aluminizing process or by codeposition. A typical pretreatment consists in chromizing before aluminizing to obtain a chromium-enriched aluminide coating. Industrially operational chromium-modified aluminide coatings include C1A (SNECMA) and PWA 32.
15
Codeposition of two (or more) elements during a single-step cementation treatment does not seem to be a widespread technique in industrial practice although it was envisaged a long time ago. Reports of combined chromium and silicon cementation of steel were published as early as 1945 [13]. Recent works based on thermodynamic calculations have demonstrated the feasibility of codepositing chromium and aluminium on nickel-base alloys [14, 16] or on iron-base alloys [15, 16], as well as aluminium and hafnium [16]. Difficulties seem to remain concerning the codeposition of aluminium and an active element such as yttrium, the solubility limit of this element being very low in nickel- or iron-based alloys. It is assumed that the available information concerning the transfer to industrial use (control of cement composition, etc.) is scarce. Incorporation of a third element by a predeposit has led to the development of platinummodified coatings. Most often platinum is electrodeposited as a layer 8/~m thick. After aluminizing, platinum is concentrated in the external region of the coating which is generally constituted of a mixture of PtA12 and fl-NiAl enriched in platinum. The addition of platinum improves
the resistance of simple fl-NiAl-based coatings, especially against type I hot corrosion. Recent work by Alprrine et al. [17] indicates that platinum may be replaced by the chemically similar palladium, which is less expensive, provided proper care is taken not to embrittle the coating during its processing
2.2. MCrAIY overlay coatings [18] MCrAIY alloys whose compositions and structures are entirely independent of the substrate can now be deposited as a result of the development of now well-established techniques such as electron beam physical vapour deposition (EB-PVD), and plasma spraying under a controlled atmosphere. As a consequence, a wide range of compositions is commercially available as shown in Table 1. 2.2.1. Structure and compositions MCrA1Y alloys are most generally multiphase materials. If the aluminium content is not too high, nickel- and cobalt-base alloys consist basically of a ductile ~ solid solution (face-centred cubic) containing a dispersion of //-NiAI or /%CoAl phase. The actual microstructure can be
TABLE 1 Examples of MCrAIY nominal compositions (wt.%) which are commercially available 119-211 Designation a
LCO ATD LCO ATD LCO ATD
Ni
29 5 5 2 7 14
Co
Fe
Bal. Bal. Bal. Bal. Bal. Bal.
Cr
AI
Y
18 18 19 23 23 30
8 11 10 12 13 5
0.5 0.3 0.5 0.3 0.6 0.5
17 22 25 25 38
6 10 12 5 11
0.5 1. 0.3 0.4 0.3
Amdry 961 Amdry 962 ATD 16 Amdry 963 ATD 1
Bal. Bal. Bal. Bal. Bal.
ATD 7 ATD 9 LN 21 Amdry 997 LN 34
Bal. Bal. Bal. Bal. Bal.
22 20 22 23 0.5
18 42 21 20 20
12 5 7.5 8.5 11
0.3 0.3 0.5 0.6 0.5
LCO 22 Amdry 995 LCO 37 Amdry 996
32 32 23 10
Bal. Bal. Bal. Bal.
21 21 30 25
7.5 8 3 7
0.5 0.5 0.5 0.6
Amdry 970 ATD 8
Bal. Bal.
24 24
8 11
Other
4 Ta 0.5 Mo
5 Ta
0.4 0.6
aATD = coatings produced by EB-PVD, Airco Temescal; LN and LCO = coatings produced by argon-shield plasma spraying, Union Carbide; Amdry = powders available for LPPS, Alloy Metals.
16
more complex as shown for example by Frances et al. [22] who have identified ~', o, M'sY and Y203 phases in addition to fl and ~ in NiCoCrAIYTa alloys. It is to be borne in mind that the microstructure of MCrAIY alloys may depend on previous heat treatments and cooling rates. If the aluminium content is sufficiently high, the fl phase percolates and contains a dispersion of ~-Ni particles. Such a structure is certainly less favourable in terms of ductility. It is to be noted that in the Ni-Cr-A1 ternary system, a phase transformation can occur at about 1000 °C whereby ~ + fl, stable at high temperature, transforms into y'-Ni3Al+ ~-Cr as represented schematically on Fig. 1 based on the phase diagrams established by Taylor and Floyd [24]. This reaction is accompanied by a significant volume variation (Fig. l(b)) which may be deleterious for the mechanical integrity of a coatingsubstrate system. It is therefore recommended that an alloy composition outside this field is selected, limiting the aluminium content, or that cobalt is added to destabilize ~' formation. The oxidation behaviour of MCrAIY-coating alloys depends on various factors (for a review see, for example, ref. 25), e.g. alloy composition, temperature, oxygen partial pressure, time and thermal cycling. A typical oxidation sequence for an MCrA1Y alloy starts with the formation of transient oxides (NiO, spinel, etc.) growing rapidly until a continuous stable oxide layer (chromia or alumina) has formed underneath. The diffusion of oxygen and metallic species being very slow in chromia and even slower in alumina, this layer ensures the oxidation resistance of the alloy. The presence of chromium in these alloys reduces the aluminium level necessary to form a protective alumina scale. The presence of as little as 5 % - 1 0 % Cr reduces
N~62Cr21AII7
(a)
..... .
(b)
Fig. 1. (a) Partial phase diagram in the N i - C r - A 1 system and (b) thermal expansion data showing the phase transformation at around 1000 °C [23].
the amount of aluminium required from 40 to about 10 at.%. Alumina scales formed on MCrA1 alloys are not adherent in thermal cycling conditions, presumably as a result of growth and thermal stresses. In order to improve the adherence of the alumina scale, so-called active elements such as yttrium, rare-earths, etc. are incorporated in these alloys in small amounts (less than 1 wt.%). Several explanations have been put forward to explain this effect but a controversy still exists relative to the precise mechanism(s) involved. Other elements (e.g. hafnium and zirconium) have been claimed to have the same effect. A recent comparative study [23] showed however that, of hafnium, zirconium and yttrium, the latter is by far the most effective addition element for improving the adherence of alumina scales formed at 1100 °C. It is to be noted that this effect is observed whether yttrium is initially present in the oxide particles or metallic phases within the coating. It seems in fact that a uniform distribution of these elements is essential. Recently, Gupta and Duvall [26] recommended the addition of both yttrium and hafnium. A major degradation (termed "hot corrosion") is the accelerated oxidation induced by condensed salts formed by pollutants present in combustion gases. Major contaminants in gas turbines are alkaline sulphates, in particular Na2SO4. The degradation modes depend on various factors: temperature range, nature of contaminants, gas composition, alloy composition [5]. Several mechanisms have been proposed (for a review, see ref. 27). At relatively high temperature (type I hot corrosion, T > 800 °C), the degradation of the scale can take place either by basic fluxing, which affects alumina-forming alloys containing less than about 15% Cr or by acidic fluxing induced by oxides such as MOO3, WO3, V205 formed by the oxidation of refractory elements. At moderate temperatures (650 °C < T < 800 °C) another form of hot corrosion (type II [28]), which affects alloys having a low chromium content, has been reported. It involves acidic fluxing of oxide scales by SO3 dissolved in molten sulphates in the presence of air. The presence of chromium is beneficial against all types of hot corrosion. It is thought that Cr203 acts as a buffer against basic fluxing mechanisms. Moreover, Cr203 is formed more rapidly than AI203. About 2 5 % - 4 0 % Cr is required for type II hot corrosion whereas 15%-20% Cr is usually
17 recommended for type I [28]. It is to be noted that too high a chromium content may have deleterious effects on the diffusional stability of the system. Silicon is particularly good in the case of type II hot corrosion (acidic fluxing) but above about 800 °C, it tends to diffuse rapidly into the substrate and may form low-melting phases or a brittle intermetallic compound. Other addition elements may play a beneficial role; tantalum, for example, although not particularly good for oxidation resistance [29], prevents the outward diffusion of potentially deleterious elements such as titanium by tying them up, with carbon, in stable MC carbides. The above considerations justify the wide use of standard NiCoCrAIY-base compositions with about 20% Co, 20% Cr, 8 % - 1 2 % AI and < 1% Y(wt.%). 2.2.2. Alternative deposition techniques for metallic coatings
Despite the flexibility they permit, the techniques commercially available for depositing MCrAIY coatings, EB-PVD and plasma spraying under inert atmosphere, remain line-of-sight processes and this, in addition to their relatively high cost, can be a real drawback for coating components having complex shapes. In view of some of the disadvantages of production techniques, alternative processes have been developed, among which electrolytic codeposition and electrophoresis seem the most promising at the moment. 2 . 2 . l 1. Electrolytic codeposition. MCrA1Y coatings have been obtained by co-electrodeposition of a dispersion of fine CrA1Y powder particles and a cobalt or nickel matrix [30]. This operation can be carried out in a barrel plating unit (Fig. 2). The barrel containing the specimens and the powder particles (10 #m diameter) is immersed in a cobalt and/or nickel conventional
Drive
\ ~ 9ear
Sol~
,n.i
level
(~
-i
Su
r
Drive shaft
Flexible
@ rotating Contact holder
I
~
(tornotor}
bath and rotated. The coated components are subsequently heat treated for alloying and diffusion bonding the coating to the substrate. This process presents the obvious economical advantages of low capital cost (claimed to be an order of magnitude less than existing commercial systems; unit processing cost about half those of competing systems [30]). Results of reliability tests do not seem to have been published as this technique has not yet been scaled up.
2.2.2.2. Electrophoresis. This process is based on the migration of fine powder particles suspended in a polar solution under a d.c. voltage. In the process developed at SNECMA [31,32], MCrA1Y alloy powders with size less than about 40 #m are deposited by electrophoresis on the components to be coated with deposition rates up to 100 #m min -~. Figure 3(a) shows the deposit with controlled porosity obtained at this stage. In order to densify them and adjust the coating composition, the components are then aluminized (vapour-phase process). The appearance of an MCrAIY coating obtained is illustrated in Fig. 3(b). With adequate electrode geometry, components with complex shapes can be uniformly coated. The facilities required are relatively simple, so the investment, as well as the running costs, is much less than for processes such as EB-PVD and plasma spraying. As shown in Fig. 3(c), these coatings confer a hot corrosion resistance comparable with that of low-pressure plasma-sprayed coatings. 2.2.2.3. Cladding. Cladding MCrA1Y sheets has been expected to produce protective coatings [33] but early experiments do not seem to have been an overwhelming success. This may have been caused by the very limited aluminium content of available sheet alloys (conventional forming processes require a low aluminium content) and possibly an inadequate yttrium distribution. It might be interesting to reconsider this technique since MCrAIY ribbons (50-100 pm thick) having more oxidation-resistant compositions can now be obtained by melt-spinning [34].
gear
2.3. Thermal barrier coatings (b
ath~e
Bearin9 /
Fig. 2. Schematic diagram of a co-electrodeposition plating unit [30]. (a) Longitudinal section; (b) transverse section.
Thermal barrier coatings (TBCs) generally consist of an oxide layer (thickness ranging from about 300 #m in turbine applications up to about 2 mm in diesel engines) plasma sprayed on top of an MCrA1Y bond layer deposited on the substrate.
18
(a) Electrophoresis + aluminizing (C3A)
1000
MCrAIY
(L.P.P.S,
1000
Cr + AI
(C1A)
Aluminide (A.P.V.S.) Uncoated
I zoo -715o
-]so Life time (h.)
Ba-
Burner rig (870°C)
(b)
(c)
Fig. 3. (a) Cross-section of an MCrA1Y coating after electrophoresis and (b) subsequent consolidation by vapour phase aluminizing, (c) hot corrosion resistance in burner rig tests, with Na2SO4 deposits, at 900°C of various coatings on IN 100 (c) [18].
The benefits of thermal barrier coatings are multiple: higher inlet gas temperatures, lower cooling flow within the metallic component, reduction of temperature transients on the metal surface. All these advantages can result in improved efficiency, extended lifetimes and simplified designs. Such ceramic coatings have been plasma sprayed on sheet metal combustor components for more than 15 years. Only recently [35] have they been used in the turbine section of commercial gas turbine engines. Thermal barrier coatings over the entire airfoil of PW4000 turbine vanes were certified in the U.S.A. for production use in 1986. This progress could be achieved by a proper control of porosity, microcrack distribution, and residual stresses in the coating. With nickel-base superalloys having reached their optimum properties, one of the present challenges t o increase gas turbine performance significantly remains the application of these coatings on highly stressed components such as turbine blades.
completely stabilized with proper additions of Y203, MgO or CaO, since ZrO2 offers a good compromise between low thermal conductivity and high thermal expansion coefficient. Yttria partially stabilized zirconia containing between 6 and 8 wt.% Y203 has been found empirically [36]
2.3. I. Zr02- (6-8)wt. % }'2 03 Figure 4 illustrates the typical aspect of the cross-section of a thermal barrier coating. Most often, the oxide is zirconia-based, partially or
Fig. 4. Typical appearance of the cross-section of a TBC showing partially stabilized zirconia (low thermal conductivity and high compressive strain tolerance), MCrA1Y bond coat (strain isolation and oxidation resistance) and superalloy substrate.
PARTIALLY STABILIZED ./ZIRCONIA
MCrAIY BOND COAT
SUPERALLOY SUBSTRATE
19 to be most resistant in thermal cycling conditions. This composition corresponds to the formation of a metastable tetragonal t' phase when the cooling rate is sufficiently high, as in plasma spraying, which is at the moment the deposition process most often used for obtaining these coatings. With proper parameter adjustment, plasma spraying gives a microcracked structure particularly strain tolerant in compression [37]. This is important as the main degradation mode (spallation) of TBCs results from fatigue cracking damage in the ceramic layer near the ceramic-bond coat interface induced by thermal strains. It should be emphasized though that the precise role of the t' phase in terms of crack propagation rate is not clearly understood and basic studies are needed to relate the fine microstructure and phase transformations of TBCs with their mechanical properties [38, 39]. Recent engine tests by Sheffler and Gupta [40] and Toriz et al. [41] have proved that TBCs obtained by EB-PVD may perform even better than plasma-sprayed coatings. This might be due to the well-adapted anisotropic microcrack network in the former case. 2.3.2. Role of the bond coat Between the oxide topcoat and substrate, an MCrA1Y bond coat confer an adequate environmental to adapt the oxide layer and the chanically.
the superalloy is required to resistance and substrate me-
2.3.2.1. Mechanical adaptation. This mechanical adaptation is more complex than a simple matching of thermal expansion coefficients. Actually the expansion coefficients of MCrAIY alloys are generally higher than for superalloys. The calculation of residual stresses in a purely elastic situation (with experimental elastic moduli) gives values well above the measured stresses. In fact, strain accommodation occurs to a large extent by plastic deformation of the bond coat as MCrAIY alloys have a very low flow stress at temperatures above 600 °C. 2.3.2.2. Environmental resistance. Thermal barrier coatings are dynamic systems in the sense that zirconia top coat, MCrA1Y bond coat and substrate superalloy interact at high temperature. Figure 5 shows, for example, the thin alumina layer formed at the zirconia-MCrA1Y interface after a 100 h heat treatment at 1100 °C. This layer has
Yttrim partial#y stabilized zirconia
Alumina
NiCrAIY
.
.
.
.
Fig. 5. Aluminalayer formedat the zirconia-MCrAlYinterface after 100h at 1100°C. formed by interaction of aluminium which has diffused from the MCrAIY alloys with oxygen present in the zirconia material. It is interesting to note that the adhesion between these layers (zirconia-alumina-MCrA1Y) is sound and that, in practice, spallation rarely occurs at the interfaces. Instead, in thermal cycling tests, failure of the TBC happens as shown on Fig. 6(a) by propagation of a main crack inside the zirconia top coat near the alumina-zirconia interface. Propagation results from compressive stresses arising in the top coat on cooling because of a thermal expansion mismatch. A more extensive oxidation of the bond coat can lead to the destruction of the barrier as shown in Fig. 6(b) as the transformation of the alloy into oxide is accompanied by a volume expansion which generates additional stresses in the zirconia layer. The evolution of the structure and composition of the MCrA1Y bond coat are likely to affect its mechanical properties and, as a consequence, to alter its ability to deform rapidly during cooling. This in turn will raise the compressive stress level in the ceramic layer, accelerating the degradation of the coating. Another degradation mode occurs by hot corrosion with liquid sulphate contaminants as illustrated in Fig. 7 [42]; when tested in cyclic oxidation at 950 °C, the weight variations are regular and no degradation occurs before 1000 cycles (test duration); however when tested in the same conditions with additional sodium sulphate contamination, a rapid degradation can be observed after only 50 cycles. Several mechanisms have been proposed to explain the accelerated degradation of the TBC in hot corrosion conditions: destabilization of zirconia, sulphidation of the bond coat, solidification of sodium sulphate in the cracks [42]. In the case of Fig. 7 these mechanisms did not seem to be operative and the
20 Yttria partially stabilized zirconia
Fig. 6. Typical degradation of a TBC: (a) propagation of a main crack parallel to the ceramic-metal interface and (b) internal oxidation of bond coat. Weight gain (mg/cm2) ~SALT
t
~/~/~ SOLIDIFICATION \
I NoContamination
01 -1 -2
: : :
\
' !
~ SALT ~ EVAPORATION~:
%.
-3
\ \
r BOND // COATING / OXIDATION
,3
o
200 400 6oo Boo Number of one hour cycles
1000
Fig. 7. Weight variations of a TBC tested at 950 °C in cyclic oxidation and in hot corrosion [42].
~MOLTENFD LIM AMAGESALI
ZIRCONIA~ TOUGHNESS ,~ REDUCTIONJ"
TEMPERATURE
Fig. 8. Schematic representation of TBC life as a function of temperature [43, 44].
explanation put forward involves a crack propagation ass~ted by the sulphate vapour.
2.3.3. Modelling A significant effort at modelling the behaviour and life of thermal barrier coatings is going on in the U.S.A., as summarized by Miller [43]. The models developed result from a predominantly phenomenological approach schematically illustrated in Fig. 8 and involve several empirical parameters. It must be emphasized that, up to now, some areas are still generally poorly described: creep and inelasticity of ceramic and bond coat, geometry effect (roughness of the metal-ceramic interface), oxidation-induced damages, sintering of oxide at high temperatures, etc. 2.3.4. Concluding remarks Thermal barrier coatings should provide the next significant increase in engine performance; Shetfler and Gupta [40] report that a zirconia
layer 250/zm thick can reduce the metal temperature by as much as 170 °C. This may be achieved through the development of more reliable coatings and of thermomechanical modelling. Moreover new ceramic compositions are required for improved coatings resistant to corrosive environments with contaminated fuels and molten salt deposits [45] and, in the long run, for coatings having a higher temperature capability than yttria partially stabilized zircouia. Finally new processes for depositing ceramic coatings will have to be developed to protect components having complex geometries.
3. Compatibility considerations Numerous considerations must be taken into account when selecting or designing a coating for a particular application; they are related to the
21
environment, the coating process, the service conditions and the substrate. Due to the drive for ever-increasing temperatures to improve engine performance and the development of advanced superalloys, considerations of coating-substrate compatibility, whether chemical or mechanical, become more and more important.
3.1. Chemical compatibility Coating and superalloy substrates generally have widely different compositions, each of them selected for different purposes: protection against oxidizing environment on the one hand and hightemperature mechanical properties on the other. The chemical potential gradients of the different elements (e.g. nickel, aluminium, chromium) across the coating-substrate interface will cause interdiffusion phenomena to occur and these effects will be all the more important when .the temperature is elevated. As a consequence, elements from the substrate may diffuse through the coating, reach the external surface and perturb the formation of the protective oxide scale. This is illustrated in Fig. 9 which represents the weight variations in cyclic oxidation at l l00°C of different superalloys (IN100, IN738 and CMSX2) coated with the same plasma-sprayed CoCrAIY alloy. The graph clearly shows that after only 250 cycles the coating deposited on IN100 and IN738 suffers extensive damage, which corresponds to internal oxidation and subsequent scale spallation, as revealed by metallographic examination, whereas on CMSX2 the coating lifetime is more than 1000 cycles. Analysis of oxidized samples suggested that, in the first two cases, the formation of the alumina scale is perturbed by diffusion of titanium from the substrates toward the external surface.
5
Am mg cn~ 2
3 2 l
Number Of cycles exposure time at HO0°C
SX2
Fig. 9. Cyclic oxidation at l l 0 0 ° C of a CoCrAIY coating low-pressure plasma-sprayed on different superalloys [46].
Conversely, interdiffusion may also affect the structure of the subcoating zone in the substrate and therefore alter or modify its mechanical properties. Several transformations can take place: (a) transformation of the initial y - y ' microstructure into y - 8 in the case of CoCrAIY/Nibase superalloys [46, 47], (b) alteration of the y' distribution, (c) modification of the distribution of grainboundary strengthening elements (carbon, etc.) [48], (d) destabilization of the reinforcing carbide fibres in directionally solidified eutectics [49], (e) formation of brittle phases which may be crack initiation sites or easy paths for crack propagation. These effects, which can be important in the case of components with thin sections, can be minimized by carefully selecting a coating composition compatible with that of the substrate. This is possible in the case of MCrAIY coatings because of the flexibility allowed by the deposition processes. Some authors [50] have suggested the interposition of a "diffusion barrier" (generally a stable ceramic material in which metallic diffusion is slow) between coating and substrate to prevent these effects. It is felt, however, that such an intermediate layer would present several drawbacks: poor thermomechanical behaviour due to thermal expansion mismatch and a preferential path for oxidation penetration once the diffusion barrier material has been reached by oxygen through a coating defect initially present or induced by local degradation in service.
3.2. Mechanical compatibility A protective coating may alter the mechanical properties of a superalloy substrate on which it is applied as a result of factors given above. It must be stressed, however, that in the long run, a coating, if adequately protective, will isolate the superalloy from the environment (oxidizing combustion gases, molten salt deposits) and will limit or suppress the corrosion-induced mechanical degradation. In a study on thin-section Ren6 80 specimens precorroded for 500 h in an engine atmosphere, Kaufman [51] has shown that the presence of a coating, whether aluminide or MCrAIY, improves the creep lifetime (at 980 and 1080 °C) by factors ranging between 2 and 6, as well as the fatigue resistance at room temperature (220-280 MPa).
22
It is generally accepted that creep properties of superalloys are not affected by a coating provided the standard ~heat treatments are respected and little interdiffusion occurs in service or during the application of the coating (Fig. 10); a possible favourable influence can occur in the long run because of the suppression of environmental degradation. The mere presence of a protective coating can significantly alter the fatigue resistance of superalloy substrates. This has been clearly shown, for example, by Veys and M6vrel [52] and by Wood and Restall [53, 54] in the case of single-crystal superalloys (Fig. 10). The reduction in fatigue lifetime is related:to early cracking in the coating with subsequent propagation inside the substrate. The fact that early cracking occurs at temperatures around 800 °C or 900 °C may be an aggravating factor for hot corrosion degradation: in this temperature range, sodium sulphate may condense o n blade surfaces and penetrate through these.cracks, abolishing the beneficial effect of the protective coating. It is important to note that the structure and composition of a protective coating change during operation a t high temperatures as a result of aluminium consumption by oxidation; the initial oR (MPa) I N -x"~- iCMSX'2unc°ated[11] •~ : CMSX-2/alurninide : CMSX-2/N CoCrA TaY 500 ! 400.
e'~e
300
200
\
\
~ ' ex
150 1'20"T(OK).10?[20 + log tr,] =25 2'9 30 :Jl Omax (MPa) o \ X~X 800
0 A
UncoatedCMSX-2 [11]
o
CMSX-2 + aluminide
¢
CMSX:2+ NiCoCrA YTa T =" 870°C = 70 HZ
" oN ~ ~ El
[3
70~ 0
.6.0p 0.1 "
1
"1'0
"Nr (cycles) . 106 100 =
Fig'. 10~ Creep and high cycle fatigue properties o f coated CMSX-2 [52].
+/~ structure of an MCrAIY coating therefore progressively transforms into ~ ( + ~ ' ) and in the end to y solid solution of nickel; this is accompanied by grain growth. No study however seems to have been published on these effects, the durations of tests performed being too short for these phenomena to be significant. Recent efforts [55-58] have concentrated on thermomechanical fatigue tests which combine mechanical strain-controlled cycling and temperature cycling; these tests offer a better simulation of the complex strain-temperature-time cycles experienced by gas engine components in service. Moreover, thermomechanical fatigue may now b e a major limiting factor in turbine vanes and blades [37], in particular for highly stressed thinwalled components. Swanson et al. [57, 58] are developing life prediction models with the emphasis on applicability to thermomechanical fatigue conditions. These models are based on the development of individual constitutive models of coatings (PWA286: NiCoCrAIYSiHf and PWA273 low activity aluminide) and PWA1480 single-crystal substrate, for which commercial production experience already exists. Results of out-of-phase tests carried out between 427 and 1038 °C indicate that at strain ranges greater than 0.5%, the overlay-coated specimens have longer lives than do those coated with the aluminide, the reverse being true at lower strain range. It is interesting to note that similar results: had been obtained by Bain [56] on a single,crystal superalloy coated with a CoCrA1Y alloy or an aluminide and tested in slightly different conditions. As emphasized by Wood [59], little work has been done to determine in depth which coating properties are of importance in controlling the mechanical behaviour of coating-substrate systems. The studies in this field have dealt mainly with the identification of degradation modes and the experimental evaluation of the incidence of coatings on a few mechanical properties of ascoated substrates. It is important to point out that coatings are found as crack initiation sites in a variety of test conditions, the cracks thus formed propagating subsequently into the substrate material and leading eventually to failure. Neither the evolution of the coating composition in service nor the presence of corrosive products are taken into account. Some attempts at modelling the coating-superalloy system have been recently published [58]. However, the lack of data relative
23 to the intrinsic mechanical properties of the coating, mainly in the case of the widespread aluminides (simple or modified), seems t o have delayed the development of an adequate description of these systems. As a consequence, the possibility of tailoring the Coating composition and/or structure to improve the mechanical properties of the coating-substrate composite does not seem to have been fully exploited.
Acknowledgments Thanks are due to m a n y people for valuable discussions and suggestions, in particular Serge Alp6rine and Pierre Jesse on palladium-modified aluminide coatings, Patrick Choquet on the oxidation of MCrA1 + X coatings, J,-M. Veys on the mechanical properties of coated systems, Ren6 Morbioli ( S N E C M A ) on electrophoretic deposits and Robert Miller (NASA) on thermal barrier coatings.
4. Conclusion Looking back on the historical evolution of high-temperature protective coatings it is interesting to note that the systems with the best inservice performance at the m o m e n t have been developed according to a largely empirical approach. Further developments can be expected along this line, for example with the introduction of alternative coating processes for depositing MCrA1Y alloys (electrophoresis, co-electrodeposition) and with the replacement of platinum by palladium in modified aluminide coatings. It has to be realized, however, that the development of more advanced processes (codeposition in single step cementation for instance) requires unavoidable preliminary theoretical treatments. More indepth investigations are needed in order to understand the degradation mechanisms better and therefore to be in a position to exploit the full potential of available coating compositions. In this respect more data and modelling are necessary concerning the intrinsic mechanical properties of coatings and also the thermomechanical behaviour of coated systems. Application of thermal barrier coatings on highly stressed components represents the next challenge for significantly improving gas turbine engine performances. This implies an active research effort in the direction of more reliable coating processes, and, depending on the applications envisaged, of new systems with higher temperature capabilities and improved hot corrosion resistance. Finally an increasing activity is noted concerning the development of high-temperature materials: titanium alloys having creep strength at temperatures above those for which oxidation resistance is adequate, intermetallic compounds, etc. Formulating new coating compositions and processes is expected to be necessary to ensure the protection of these materials.
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47 P. Mazars, D. Manesse and C. Lopvet, Interdiffusion of MCrA1Y coatings with substrates, in High Temperature Alloys for Gas Turbines and Other Applications 1986, Reidei, Dordrecht, 1986, pp. 1183-1192. 48 A. Strang, E. Lang and R. Pichoir, AGARD SMP, CP356, 11/1-11/35, 1983. 49 J.-M. Veys and R. M6vrel, Mater. Sci. Eng., 88 (1988) 253-260. 50 J. P. Coad, D. S. Rickerby and B. C. Oberlander, Mater. Sci. Eng., 74 (1985) 93-103. 51 M. Kaufman, Examination of the influence of coatings of thin superaUoy sections, NASA Contract. Rep. CR-1215 (1972). 52 J.-M. Veys and R. M6vrel, ASM Technical Conference, Paris, September, 1987, American Society for Metals, Metals Park, OH. 53 M. I, Wood and J. E. Restall, in High Temperature Alloys for Gas Turbines and Other Applications, Liege, 1986, pp. 1215-1226. 54 M. I. Wood, Proc. 1st ASM Europe Technical Conference on Advanced Materials and Processing Techniquesfor Structural Applications, Paris, September 1987, ONERA, 1987, pp. 179-188. 55 T. E. Strangman, PhD Thesis, University of Connecticut, 1978. 56 K. R. Bain, AIAA-85-1366, 1985, pp. 1-6. 57 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, NASA Contract. Rep. CR 174952 (1986). 58 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, NASA Contract. Rep. CR 179594 (1987). 59 M. I. Wood, Mater. Sci. Eng., A120/121 (2) (1989) 633.