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Coatings for gas turbine materials and long term stability issues
Materials & Design Materials and Design 26 (2005) 223–231 www.elsevier.com/locate/matdes
Coatings for gas turbine materials and long term stability issues M.J. Pomeroy
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Department Materials Science and Technology and Materials and Surface Science Institute, University of Limerick, Ireland Available online 7 June 2004
Abstract This paper reviews protective coatings against the high temperature oxidation and corrosion of gas turbine components. Having briefly reviewed the development of gas turbine materials over the past 50 years, the need for corrosion protective coatings and their routes of application and chemistries are explored. The effects of varying coating chemistries and application methods is examined in the context of the major corrosive degradation mechanisms which operate in aircraft and industrial gas turbines. A case study relating to the interdiffusion of coatings and a typical third generation Ni-based alloy is presented which shows that this phenomenon may be of importance with respect to coating life. Finally the paper briefly investigates thermal barrier coatings and how their failure is attributable to the oxidation of the bond coats to which they are attached. Ó 2004 Elsevier Ltd. All rights reserved. Keywords: Gas turbine materials; Aluminide coatings; MCrAlY coatings; Interdiffusion; TBCs
1. Alloy development Since the use of stainless steels for the blades and vanes of the first gas turbine engines, much alloy development has taken place, both with respect to composition and processing [1]. Fig. 1 shows how nickelbased alloy compositions have changed since 1965 and clearly shows that the chromium content has been drastically reduced from about 15 wt% to about 3 wt% and that aluminium contents have increased to about 5 wt%. A key indication from Fig. 1 is the increase in refractory element content, such that over the 30 years period between 1965 and 1995, additions of tantalum, rhenium, tungsten and molybdenum have increased from 8 to 20 wt%. The reasons for these changes in composition are two fold and based on the strengthening mechanisms involved in developing Ni-based alloys for optimum creep resistance at temperatures up to 1100 °C. A typical microstructure of a third generation single crystal Ni-based superalloy is shown in Fig. 2 and clearly shows that the alloy comprises a series of c0 cuboidal precipitates with an interprecipitate ‘‘cement’’ of
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0261-3069/$ - see front matter Ó 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2004.02.005
a c Ni-matrix. The composition of c0 is effectively Ni3 Al and thus the c matrix contains the refractory alloying elements Ta, Re, W and Mo, although much Ta enters c0 due to its solubility in this phase. The presence of these refractory elements in solid solution in the matrix gives rise to significant solid solution strengthening due to the strength of bonding to Ni [1]. Furthermore, the diffusion effects required for dislocation climb aided creep or diffusion creep are also slowed by the strong interatomic bonding between Ni and the refractory elements [1]. Clearly, as the refractory element content is increased, then alloys become stronger (due to solid solution strengthening) and more creep resistant (due to stronger interatomic bonding). Thus, as shown in Fig. 1 optimised c0 and therefore Al contents were arrived at in the early 1980s and further improvements have only been realised via solid solution strengthening additions. Processing techniques have also led to improved alloy strength and creep resistance with high purity alloys being prepared via vacuum melting and casting techniques. Furthermore, since creep failure is almost always a grain boundary initiated phenomenon due to either grain boundary sliding or vacancy condensation on boundaries perpendicular to the applied tensile stress, the development of directionally solidified or single crystal turbine components has resulted in further creep
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applied to decrease the temperature of the hottest parts of the turbine components by up to 170 °C.
25
wt. % element
20
2. Oxidation and corrosion degradation processes 15
10
5
0 1960 Cr
1970 Ti
1980 year
1990
2000
Al
Mo+Ta+W+Re
Fig. 1. Variations in alloying element additions to Ni-based alloys with year of alloy introduction (based on data in [2]).
In order to understand the development of corrosion protective coatings, it is necessary to appreciate the processes by which oxidation and corrosion occur and how the mechanisms by which they occur depend on environment and temperature. Three accelerated degradative processes occur depending on temperature and these can be defined in order of increasing temperature as: Type II hot corrosion, Type I hot corrosion and oxidation (see Fig. 3). Type II hot corrosion occurs at temperatures in the range 600–850 °C and involves the formation of base metal (nickel or cobalt) sulphates which require a certain partial pressure of sulphur trioxide for their stabilisation. These sulphates react with alkali metal sulphates to form low melting point compounds which prevent a protective oxide forming. Indeed, the oxide formed is a striated one as demonstrated in the work of Viswanathan [3]. As discussed later, Nicholls et al. [4] have shown that Ni–Cr–Al materials resist this form of corrosion most effectively. Type I hot corrosion involves the transport of sulphur from a sulphatic deposit (generally Na2 SO4 ) across a preformed oxide into the metallic material with the formation of the most stable sulphides. Once stable sulphide formers (e.g. Cr) are fully reacted with the sulphur moving across the scale, then base metal sulphides can form with catastrophic consequences as they are molten at the temperatures at which Type I hot corrosion is observed (750–950 °C) [3,5]. Thus, the formation of NiS2 (molten at 645 °C) and Cox Sy (lowest
property improvements. Thus, directionally solidified components are cast such that there are a minimum of grain boundaries perpendicular to the major tensile stress and single crystal components, of course, contain no grain boundaries. All of this alloy development has been made to improve mechanical properties to facilitate higher component temperatures. However, the low chromium and aluminium contents mean that these materials do not have the necessary intrinsic resistance to oxidation and corrosion required for the long term operation associated with gas turbine operation. Accordingly, coatings must be applied to gas turbine components for oxidation or oxidation/corrosion resistance. In addition, due to the demand to increase turbine inlet temperatures and thus cycle efficiencies, ceramic insulating coatings can be
corrosion rate (arb. units)
Fig. 2. Typical microstructure of a Ni-based superalloy showing cuboidal c0 and c matrix (bar ¼ 1 lm).
500
700 900 temperature (˚C)
Type II
Type I
1100
Oxidation
Fig. 3. Schematic representation of rate-temperature curves for Type II hot corrosion, Type I hot corrosion and oxidation (alumina former).
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log (parabolic rate constant)
liquidus 840 °C) can cause degradation levels which are serious enough to cause major component degradation. The most suitable materials which can resist Type I hot corrosion are PtAl2 –(Ni–Pt–Al) coatings and M–CrAlY coatings containing up to 25 wt% Cr and 6 wt% Al [6]. The oxidation of metals depends on the rates at which anion or cation transport can occur through the crystal lattice or along grain boundaries in the oxide [7]. In alloys, which oxide is the most stable depends on the oxide dissociation pressure which is lowest for Al and Cr compared to the typical base elements Fe, Co and Ni. For Ni–Cr alloys containing more than about 10 at.% Cr, a continuous protective chromia layer is formed. A higher Cr level (25 at.%) is required for Co-based alloys because Cr diffuses more slowly in Co and so cannot form a continuous chromia layer at lower concentrations [7]. Chromium oxide can itself oxidise at temperatures greater than about 850 °C to a volatile CrO3 compound. Because of this, the use of aluminium additions for oxidation resistance is preferred at this temperature and above for key components such as those used in gas turbines. There is added advantage, as shown schematically in Fig. 4, because the rate of formation of aluminium oxide is slower than for chromium oxide at the same temperature. Addition levels of Al to Ni for the formation of a continuous alumina scale are as for chromia formation (10 at.%). Alumina formation at lower Al contents can be induced by mixed Cr + Al additions (see Fig. 4) since Cr ‘‘getters’’ oxygen allowing alumina to form at lower aluminium activities (contents) [7]. The aluminium and chromium contents referred to above apply to isothermal oxidation conditions. However, when thermal cycling conditions prevail, oxide scales can spall from the substrate surface due to thermally induced stresses [8]. The oxidation resistance under such conditions can be markedly improved by the addition of so-called reactive elements (Y, Hf, Ce) to
Cr
Al + Cr Al reciprocal temperature
Fig. 4. Schematic Arrhenius plot for oxidation of alloys containing Cr, Cr + Al and Al.
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alloys and coatings. An exhaustive search for the reasons behind this effect has finally shown that yttrium and rare earth metals segregate to grain boundaries within alumina scales causing a reduction in Al and O transport rates through the oxide and thus reduced oxidation rates [9]. Furthermore, reactive elements combine with sulphur and phosphorus impurities in metallic materials and coatings with the result that these impurities cannot selectively diffuse to the surface and contaminate the oxide–metallic interface. This gives extremely good scale adherence and categorically explains why minor yttrium or cerium (<0.8%) additions to oxidation resistant coatings and alloys greatly improve oxidation resistance.
3. Corrosion resistant coatings-basic processes From the above analysis, aluminium or chromium rich coatings are required for oxidation and corrosion resistance with the former type being required for resistance to oxidation and Type I hot corrosion and the latter type for Type II hot corrosion. Two categories of coatings arise: diffusion coatings and overlay coatings. Diffusion coatings are formed by a chemical vapour deposition type processes e.g. aluminising or chromising or indeed codeposition of both Al and Cr. Whilst these processes are primarily CVD-based, they are more commonly referred to as diffusion coatings since their application involves interdiffusion between deposited Al or Cr and the substrate onto which they are coated. Thus, for example the aluminising of a nickel-based superalloy gives rise to a coating of generic composition NiAl (b-nickel aluminide). Several variants of these ‘‘ising’’ processes arise. These are: pack [10], out of pack [10], CVD [11] and fluidised bed CVD [12]. With pack aluminising or chromising components to be coated are placed in a ‘‘pack’’ of a diluent, typically alumina, a halide activator (e.g. ammonium fluoride) and metal powder e.g. Al or Cr and heat treated in argon. For the out of pack process, the component is located downstream from the ‘‘pack’’ such that the metal halide gas produced by the pack impinges on the substrate and the metal is deposited on the surface where it undergoes interdiffusion with the substrate. In the CVD process Al and HCl are reacted under controlled conditions and passed under a positive pressure into the coating part of the system where they enable aluminide formation [13]. A particular benefit of this system is the cleanliness of the coatings yielding improved oxidation resistance [11] and the fact that the gas can be passed through, as well as over, turbine components with cooling holes such that internal and external surfaces are coated as well as the cooling holes. In contrast, pack aluminising can cause cooling holes to be blocked with diluent particles which require removal after processing. An additional
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benefit of the CVD process is that the chemistry can be carefully controlled enabling additional elements to be incorporated into coatings at precise levels [13]. Such ideas form the basis of an EU R&D Growth project (G4RD-CT-2000-00319). The fluidised bed aluminising process [12] involves the suspension of the component in a bed of inert particles (e.g. alumina) by a gas stream comprising the gases necessary for aluminising in a carrier gas. Given that fluidised bed technology is characterised by excellent solids mixing and heat transfer but poor gas mixing, temperature control may be excellent but inlet gas chemistries must be carefully controlled. Overlay coating involves the application of coatings to substrates using physical deposition processes. Typical application methods include: thermal spraying, physical vapour deposition (PVD) electron beam physical vapour deposition (EBPVD), ion plating/sputtering and electroplating [14].
4. Corrosion resistant coatings-types and chemistries 4.1. Overlay coatings Overlay coatings typically comprise b þ c0 aluminide in a c matrix [1] and are of typical composition (Ni, Co)15–28 wt% Cr, 4–18 wt% Al, 0.5–0.8 wt% Y [6]. The relative amounts of Ni and Co depend upon: (a) coating ductility requirements (>Ni) and (b) corrosion resistance (>Co) [6]. The excellent oxidation and corrosion resistance of the coatings is afforded by the formation of highly adherent alumina scales which grow slowly as a result of the reactive element effect attributable to yttrium. In addition, their high Cr contents make them useful protective coatings against Type II hot corrosion. These overlay coatings are deposited using thermal spray (Ar shrouded plasma (APS) or low pressure plasma (LPPS)) techniques or electron beam physical vapour deposition (EBPVD). The EBPVD technique is typically preferred for high quality coatings since a certain amount of oxidation of the coating particles typically occurs during thermal spraying giving rise to nanoscale oxide particles at splat boundaries. However, the thermal spray processes are frequently used in practice. Overlay coatings show interdiffusion effects with alloy substrates and Itoh and Tamura [15] indicate the rate of interdiffusion decreases in the order NiCrAlY > CoCrAlY > NiCoCrAlY > CoNiAlY, where NiCo represents higher Co contents compared to Ni and CoNi higher Co contents. These results, collected using vacuum plasma sprayed coatings were independent of the three substrates used in the experimental programme. The most recent technological advance in overlay coating technology is the development of so-called smart coatings by Nicholls et al. [4]. These coatings attempt to
address the problems associated with the differences in temperature over the surface of an airfoil. Temperatures vary from a maximum of the order of 1100 °C at leading and trailing edges to about 650 °C in the centre of the airfoil surfaces and near airfoil roots. Because of this, the nature of environmental degradation varies from oxidation through Type I hot corrosion to Type II hot corrosion as indicated in Fig. 2. A practical example of this variation in corrosion type with temperature is graphically portrayed by Viswanathan [5] in his analysis of the failure of airfoils. Following extensive laboratory studies employing alkali–metal sulphates in simulated gaseous environments, Nicholls et al. [4] have been able to very carefully define optimised chromium and aluminium contents for oxidation and corrosion protection in the temperature ranges referred to above. As might be expected, high Cr contents (>40 wt%) and low Al contents (6–8 wt%) are most suitable for protection from Type II hot corrosion (650 °C). For protection from Type I hot corrosion (800–950 °C), roughly equal amounts of Cr and Al represent optimised compositions. Oxidation protection (1100 °C), is best afforded by coatings containing 25 wt% Cr and 14 wt% Al. The final coating developed by Nicholls et al. [4] comprises a commercial base coating (Co–32Ni–21Cr–8Al–0.5Y) adjacent to the substrate, a Cr enriched layer of variable composition from Ni–60Cr–20Al to Ni–35Cr–40Al and a surface layer of composition Ni–15Cr–32Al. These multi-layer coatings have been shown to out perform typical a commercial Pt modified aluminide and an Al enriched version of the base coat at 700–800 °C. Whilst this technology appears a step forward, issues related to coating ductility and additional improvements afforded by cobalt with respect to corrosion resistance, may be of additional importance. 4.2. Diffusion coatings 4.2.1. Chromium rich coatings Effectively, chromising results in Cr enrichment of the surface layers which results in either greater Cr levels in the bulk element solid solution (i.e. Fe, Co, Ni) or supersaturated solid solutions plus aCr phases [16]. As for aluminising, chromising is effected by heating components in a gas rich in a volatile Cr compound (e.g. CrCl3 ). Chromising is of particular use for resistance toward Type II hot corrosion and has also been found to be of major benefit in protecting Ni-based alloys from corrosion by sulphatic deposits in chemical plant. That said, Ni–50Cr alloys have been shown to be equally as corrosion resistant. 4.2.2. Aluminide coatings Aluminising is carried out by two different processes which differ in terms of the activity of aluminium in the gas phase and the temperature at which the aluminising
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process is carried out [10]. The low activity high temperature (LAHT) is a ‘‘one step’’ process as far as the development of a b-NiAl coating is concerned. On the other hand the high activity low temperature route typically gives rise to a d-Ni2 Al3 coating which requires subsequent heat treatment to convert it to b-NiAl. As might be expected, the mechanisms by which the coatings grow is also different. For the LAHT process the aluminium activity is insufficient for it to be the predominant diffusing species and accordingly, coatings form by the diffusion of nickel from the alloy substrate into the region of the coating. For the HALT process the aluminium activity in the gas phase and at the surface of the coating is high enough to facilitate the inward diffusion of Al into the alloy substrate. The differing growth modes may be of importance when considering coating integrity, since coatings which grow by outward Ni diffusion can trap diluent particles (alumina) in aluminising packs within them. Which coating process is selected depends upon a number of features e.g. heat treatment specifications for the substrate alloy, nature of packs available, integrity issues, etc. [10]. 4.2.3. Platinum modified aluminide coatings Platinum modified nickel aluminide coatings exist in two forms, a two phase PtAl2 – (Ni–Pt–Al) (Fig. 5(a)) or single phase Pt modified b-NiAl (Fig. 5(b)). The manner by which the coatings are formed involves an initial deposition of a layer of platinum, typically 6 lm thick either by electrodeposition or ion plating. Following a post coating annealing process to diffusion bond the platinum to the substrate, aluminising gives rise to the platinum modified NiAl coating. If the prealuminising heat treatment is carried out such that significant platinum diffusion into the surface layers of the substrate occurs, the aluminide formed is of the single phase variety. If the pre-aluminising anneal is conducted such that platinum diffusion effects are small, then the bi-phasic PtAl2 – (Ni–Pt–Al) forms [17]. One of the original reasons for platinum modification was to try to suppress the diffusion of refractory elements (W, Mo) into the surface layers of aluminide coatings. Results obtained by Leyens et al. [18] clearly indicate that this effect is not wholly realised and the subsequent work by Zhang et al. [19,20] has shown that the most probable beneficial effect of platinum in aluminide coatings is to suppress void formation at the coating-alumina scale interface. This greatly improves scale adhesion and accordingly oxidation and corrosion resistance. Two clear examples of the improvement of the enhanced oxidation resistance afforded by Pt modification of aluminide coatings are reported by Krishna et al. [17] and Purvis and Warnes [21]. Results from both of these studies show improved resistance to alumina scale spallation as Pt contents are increased.
Fig. 5. Microstructures of Pt modified aluminide coatings on Ni-based superalloy: (a) two phase PtAl2 + (Ni–Pt–Al) coating; (b) single phase (Ni–Pt–Al) coating (bar ¼ 20 lm).
5. Coating-substrate interdiffusion with diffusion aluminide coatings Coatings have a totally different composition from the alloys to which they are applied. For example a stoichiometric NiAl coating contains 31.5 wt% Al and 68.5 wt% Ni. In contrast, as is apparent from Fig. 1, typical Al contents of Ni-based superalloys are 5 wt% and Ni contents range between 60 and 75 wt%. In order to more clearly illustrate this point, Fig. 6 shows concentration profiles for Ni, Pt and Al across the coating and diffusion zones and into the alloy substrate for a two phase Pt modified NiAl coating on a Ni-based superalloy. Because of these concentration differences, coating – substrate interdiffusion effects arise when alloy-coating diffusion couples are put into service. Accordingly, Ni is expected to diffuse from the substrate into the coating and Pt and Al are expected to diffuse countercurrently. Basuki et al. [22] and Chen and Little [23] have studied the effects of such interdiffusion effects on microstructure and observed the degradation of the as coated single phase b-NiAl to b þ c0 (Ni3 Al). In addition, the
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elemental concentration (at.%)
80 70 60 50 40 30 20 10 0 0
50 100 distance from coating surface (µm)
Fig. 6. Compositional variations for Ni(r), Pt(d) and Al(j) across coating, diffusion layer and into alloy for cross-section shown in Fig. 5(a) .
interdiffusion effects caused microstructural change in the substrates. In particular, Chen and Little [23], who used a Pt modified NiAl coating concluded that Pt diffusion into the substrate resulted in the precipitation of topologically close packed (TCP) phases. These phases are known to be deleterious to the mechanical behaviour of Ni-based superalloys and much work has gone into predicting compositional and temperature ranges over which they form such that these compositional and temperature ranges can be avoided [1]. Recent work at the University of Limerick has looked at the interdiffusion of Pt modified aluminide coatings with CMSX-4 and CMSX-10 alloy substrates. Fig. 7 shows typical microstructures of the coatings on the latter substrate after isothermal exposure in air for 750 h at 950 °C and 375 h at 1100 °C. After both exposure periods, the PtAl2 + (Ni–Pt–Al) coating has transformed to a single phase (Ni–Pt–Al) coating and bright grain like features arise in the coating. These light grains have a nickel rich composition and contain tantalum. Since Ta is soluble in c0 (Ni3 Al) but not in b-NiAl, the original Pt modified b-NiAl has transformed to a mixed b þ c0 . This is expected on the basis of the data given in Fig. 8 which presents data for variations in Ni, Pt and Al contents of the exterior layer of the coatings with exposure time. Given that some 36 at.% Al is the minimum required for b to remain stable, it can be seen from this figure that after 375 h at 1100 °C and 750 h at 950 °C the b phase must convert to b þ c0 . There is still however, sufficient Al in the coating to confer oxidation protection. Although this is the case, the Ta rich c0 grains are prone to catastrophic oxidation and thus if the alumina scale formed above them fails then pitting oxidation may arise. Examples of this have been observed after
Fig. 7. Microstructural changes in coating shown in Fig. 4(a) after oxidation exposure at 950 °C for 750 h (a) and 1100 °C for 375 h (b) (bar ¼ 20 lm).
1500 h exposure at 1100 °C. Similar effects have been observed for the CMSX-4 alloy. However, in contrast to the results of Chen and Little [23], this work has shown no evidence of Pt diffusion into the substrate. The results obtained do however complement the results presented by Chen and Little in that Pt concentration gradients rapidly homogenise as suggested by Fig. 8. Close scrutiny of Fig. 7 shows the occurrence of needle like phases in the substrate alloy. These phases are Re, W and Cr rich and X-ray diffraction shows them to be the TCP phase referred to as r-phase. Fig. 9 plots the depth of penetration of these needle like precipitates into the alloy with respect to the square root of exposure time. Two factors are apparent: (a) an incubation period exists at 950 °C such that penetration after 188 h is only slightly greater than in as coated specimens and (b) straight line relationships arise which are indicative of a diffusion controlled microstructural change mechanism. Detailed examination of the alloy microstructure at the root of a needle-like precipitate, shows that the needle grows along a pre-existing c matrix channel. In addition,
concentration (atomic percent)
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78 65 52 39 26 13 0 0
200 400 600 time of isothermal exposure (hours) Pt 950˚C
Pt 1100˚C
Al 950˚C
Al 1100˚C
Ni 950˚C
Ni 1100˚C
800
Fig. 8. Compositional changes in exterior 16 lm of coating after oxidation exposure at 950 and 1100 °C.
depth of penetration of needle phases into substrate (µm)
250
200
150
100
50
0 0
500 1000 1500 2000 square root of isothermal exposure time (s1/2 ) 950˚C
1100˚C
Fig. 9. Effect of exposure time on depth of penetration of needle like phases into CMSX-10 alloy at 950 and 1100 °C.
in contrast to an original alloy microstructure comprising cuboidal c0 precipitates of highly regular shape and size (see Fig. 2), cuboids adjacent to the needle like precipitates appear to have combined to form an envelope around the needle. Away from this envelope, other cuboids have coalesced to form oblong shaped grains a microstructural feature referred to as rafting [24]. Sims et al. [1] draw attention to decreases in creep rupture and impact strength properties associated with TCP phase formation and Feng et al. [25] have shown that rafted microstructures compromise creep rupture properties. Accordingly, the coating-substrate interdiffusion effects shown for the CMSX-10 alloy which are also observed for the CMSX-4 alloy may give rise to a weakening of the substrate superalloy with increasing exposure time (see Fig. 10).
Fig. 10. Detail of development of needle like phase and microstructural rafting in CMSX-10 alloy (bar ¼ 5 lm).
Additional detailed concentration profile determinations after different times of isothermal exposure at 950 and 1100 °C for both CMSX alloys clearly show that Ni diffuses from the substrate into the coating and Al diffuses from the coating into the substrate and that the depth of the diffusion effects into the alloy increases with time. This data, in addition to the microstructure adjacent to the needle like phases, leads to the conclusion that Al diffuses into the bulk of the alloy via c matrix channels and Ni diffuses outward along the same route. Both of these effects lower the Ni activity in the c matrix channels with the result that the activities of solid solution strengthening elements (Re, W, Cr) increase to levels where they become insoluble as explained by Blavette et al. [26]. This insolubility leads to their precipitation at the route of growing r-needles. In areas away from the growing needles, decreases in Ni activity and increases in Al activity within c matrix channels probably approach the c0 (Ni3 Al) stoichiometry. If this occurs then the c channel will be converted into c0 giving rise to what appears to be cuboid coalescence and hence rafting. These interdiffusional effects are clearly long term factors and accordingly may influence the long term stability of both coating and substrate with respect to the functions they are expected to perform. The possibility of decreases in oxidation resistance of the coatings and decreases in mechanical properties of the alloy component may give some cause for concern. That said, there is little evidence of failures due to such interdiffusional effects although long term experience has only been gained with second generation turbine materials. The observations reported above may, however, indicate that further development of aluminide coatings is needed for third generation gas turbine alloys such as CMSX-10 if interdiffusional effects are not to be life limiting.
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6. Thermal barrier coatings The application of 0.1–0.2 mm thick thermally sprayed or EBPVD coatings of partially stabilised zirconia result in decreases in metal surface temperatures of some 170 °C or, for the same metal surface temperatures, increases in turbine inlet temperatures of 170 °C [27]. Brindley and Miller [27] indicate that such increases in turbine inlet temperatures can result in increases in thrust of 5% and efficiency increases of the order of 1% with attendant fuel economies. Partial stabilisation of the zirconia is achieved by additions of yttria and 6–8 wt% levels appear optimum. Whilst thermally sprayed coatings are adequate for stationary components (e.g. nozzle guide vanes), EBPVD coatings are preferred for rotating components. Examples of the structures of these coatings can be found in [28]. The columnar structure of the EBPVD coating facilitates a certain degree of strain tolerance and this is why this coating type are preferred on rotating parts. Thermal barrier coatings (TBCs) are applied to bond coats which were initially simple NiAl coatings. MCrAlY or Pt modified NiAl coatings are now the norm. Failure processes associated with TBCs are invariably associated with their decohesion from the alumina scale which grows on the bond coat. Originally, Al depletion in the bond coatings gave rise to the formation of spinels which due to their faster growth rate and voluminous oxide caused TBCs to spall [29]. A recent paper [30] also draws attention to this failure mode. It is now however becoming apparent that TBC failure is associated with the growth and adhesion properties of the alumina scales formed on MCrAlY and Pt modified NiAl bond coats [31–33]. Clarke et al. [34] and Tolpygo et al. [35] have shown how stresses develop during the growth of alumina and the fact that the surface of alumina scales are rumpled due to enhanced grain boundary growth rates. The stress build ups as well as oxide rumpling cause alumina to decohere from the bond coat as faster growing grain boundary regions of the oxide effectively push the TBC away from flatter regions of the oxide. Due to crack initiation at the TBC-bond coat interface, TBC decohesion occurs. It would seem that the work completed by Kim et al. [36] shows that the type of thermal cycling testing, and therefore service condition, is important with respect to TBC life. This work shows that Pt modified bond coats behave in a superior manner to MCrAlY coats if cycle times are long (>1 h) whereas the opposite is true for short cycle times (10 min).
7. Conclusion From the above discussion of coatings for gas turbines, it is clear that many types arise and which coating is best depends upon the operating conditions in which
it is to be employed. For Types II hot corrosion MCrAlY coatings present the best protection method whilst for Type I hot corrosion Pt rich coatings and MCrAlY coatings appear best. Generally, for oxidation protection and TBC bond coats Pt modified NiAl would appear best particularly if the aluminising is performed using gas CVD aluminising. It does however become quite clear that alloy-coating interdiffusion reactions have received little attention in the scientific literature, thus coatings are typically only tested with respect to their corrosion and oxidation resistance. Accordingly, whilst the current coating-turbine alloy systems are adequate, the drive for aero-turbine CO2 emission reduction by the EU and the improved economics of ultrahigh efficiency combined cycle electricity generation will see the introduction of alloys with high refractory element contents to meet the exacting requirements. It may be that, for these alloys, coating substrate interactions may control life expectancy and accordingly this area of research will become of significant importance.
Acknowledgements The author wishes to thank Mr Michael Reid who is conducting his PhD programme on Ni-based alloyaluminide coating interdiffusion effects for his help in preparing this manuscript. The coating work at the University of Limerick is sponsored by the Materials and Surface Science Institute under the auspices of the Higher Education Authority’s Programme for Research in Third Level Institutions 1999–2002. Thanks are also due to SIFCO Turbine Components Ltd., Carrigtwohill, Co. Cork for their assistance with Mr Reid’s project.
References [1] Sims CT, Stoloff NS, Hagel WC. Superalloys II. New York: Wiley; 1987. [2] Caron P, Khan T. Evolution of Ni-based superalloys for single crystal gas turbine blade applications. Aerospace Sci Technol 1999;3(8):513–23. [3] Viswanathan R. An investigation of blade failures in combustion turbines. Eng Fail Anal 2001;8(5):493–511. [4] Nicholls JR, Simms NJ, Chan W, Evans HE. Smart overlay coatings-concept and practice. Surf Coat Technol 2002;149(23):236–44. [5] Eliaz N, Shemesh G, Latanision RM. Hot corrosion in gas turbine components. Eng Fail Anal 2002;9(1):31–43. [6] Goward GW. Progress in coatings for gas turbine airfoils. Surf Coat Technol 1998;109(1–3):73–9. [7] Wood GC, Stott FH. Oxidation of alloys. Mater Sci Technol 1987;3(7):519–30. [8] Mevrel R. Cyclic oxidation of high-temperature alloys. Mater Sci Technol 1987;3(7):531–5. [9] Pint BA. Experimental observations in support of the dynamic – segregation theory to explain the reactive–element effect. Oxid Met 1996;45(1–2):1–37.
M.J. Pomeroy / Materials and Design 26 (2005) 223–231 [10] Pichoir R, Hauser JM. Current status of coatings for high temperature applications. In: Environmental degradation of high temperature materials. Inst. Metall. Spring Residential Conf., Series 3, Number 13, Vol. 1; 1980. p. 6/1–6/21. [11] Warnes BM, Punola DC. Clean diffusion coatings by chemical vapor deposition. Surf Coat Technol 1997;94-95(1-3):1–6. [12] Voudouris N, Christoglou C, Angelopoulos GN. Formation of aluminide coatings on nickel by a fluidised bed CVD process. Surf Coat Technol 2001;141(2-3):275–82. [13] Squillace A, Bonetti R, Archer NJ, Yeatman JA. The control of the composition and structure of aluminide layers formed by vapour aluminising. Surf Coat Technol 1999;120–121:118–23. [14] Stringer J. Role of coatings in energy producing systems: an overview. Mater Sci Eng 1987;87:1–10. [15] Itoh Y, Tamura M. Reaction diffusion behaviors for interface between Ni-based super alloys and vacuum plasma sprayed MCrAlY coatings. J Eng Gas Turb Power 1999;121(3):476–83. [16] Castle AR, Gabe DR. Chromium diffusion coatings. Int Mater Rev 1999;44(2):37–58. [17] Krishna GR, Das DK, Singh V, Joshi SV. Role of Pt content in the microstructural development and oxidation performance of Pt-aluminide coatings produced using a high-activity aluminizing process. Mater Sci Eng A 1998;251(1–2):40–7. [18] Leyens C, Pint BA, Wright IG. Effect of composition on the oxidation and hot corrosion resistance of NiAl doped with precious metals. Surf Coat Technol 2000;133:15–22. [19] Zhang Y, Haynes JA, Lee WY, Wright IG, Pint BA, Cooley KM, et al. Synthesis and cyclic oxidation behavior of a (Ni, Pt) Al coating on a desulfurized Ni-base superalloy. Metall Mater Trans A 1999;30(10):2679–87. [20] Zhang Y, Haynes JA, Lee WY, Wright IG, Pint BA, Cooley KM, et al. Effects of Pt incorporation on the isothermal oxidation behavior of chemical vapor deposition aluminide coatings. Metall Mater Trans A 2001;32(7):1727–41. [21] Purvis AL, Warnes BM. The effects of platinum concentration on the oxidation resistance of superalloys coated with single-phase platinum aluminide. Surf Coat Technol 2001;146:1–6. [22] Basuki E, Crosky A, Gleeson B. Interdiffusion behaviour in aluminide-coated Rene 80H at 1150 °C. Mater Sci Eng A 1997;224(1-2):27–32. [23] Chen JH, Little JA. Degradation of the platinum aluminide coating on CMSX4 at 1100 °C. Surf Coat Technol 1997;92(1– 2):69–77.
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[24] Bressers J, Arrell DJ, Ostolaza K, Valles JL. Effect of an aluminide coating on precipitate rafting in superalloys. Mater Sci Eng A 1996;220(1-2):147–54. [25] Feng H, Biermann H, Mughrabi H. Microstructure-based 3D finite element modelling of lattice misfit and long-range internal stresses in creep-deformed nickel-base superalloy single crystals. Mater Sci Eng A 1996;214:1–16. [26] Blavette D, Cadel E, Deconihout B. The role of the atom probe in the study of nickel-based superalloys. Mater Charact 2000;44:133– 7. [27] Brindley WJ, Miller RA. TBCs for better engine efficiency. Adv Mater Process 1989;8:29–36. [28] Miller RA. Thermal barrier coatings for aicraft engines – history and directions, in: Proceedings of the Thermal Barrier Coating Workshop. NASA as NASA Conference Publication 3312; 1935. p. 17–34. [29] Sohn YH, Biederman RR, Sisson Jr RD. Microstructural development in physical vapour-deposited partially stabilized zirconia thermal barrier coatings. Thin Solid Films 1994;250(12):1–7. [30] Shillington EAG, Clarke DR. Spalling failure of a thermal barrier coating associated with aluminum depletion in the bond-coat. Acta Mater 1999;47(4):1297–305, MAR 10 1999. [31] Tolpygo VK, Clarke DR, Murphy KS. Oxidation-induced failure of EB-PVD thermal barrier coatings. Surf Coat Technol 2001; 146:124–31. [32] Evans AG, Mumm DR, Hutchinson JW, Meier GH, Pettit FS. Mechanisms controlling the durability of thermal barrier coatings. Prog Mater Sci 2001;46(5):505–53. [33] Yanar NM, Meier GH, Pettit FS. The influence of platinum on the failure of EBPVD YSZ TBCs on NiCoCrAlY bond coats. Scr Mater 2002;46(4):325–30. [34] Clarke DR, Christensen RJ, Tolpygo V. The evolution of oxidation stresses in zirconia thermal barrier coated superalloy leading to spalling failure. Surf Coat Technol 1997;94-95(1-3):89– 93. [35] Tolpygo VK, Clarke DR. Surface rumpling of a (Ni, Pt) Al bond coat induced by cyclic oxidation. Acta Mater 2000;48(13):3283– 93. [36] Kim GM, Yanar NM, Hewitt EN, Pettit FS, Meier GH. The effect of the type of thermal exposure on the durability of thermal barrier coatings. Scr Mater 2002; 46(7):489–95.
Coatings for gas turbine materials and long term stability issues M.J. Pomeroy
*
Department Materials Science and Technology and Materials and Surface Science Institute, University of Limerick, Ireland Available online 7 June 2004
Abstract This paper reviews protective coatings against the high temperature oxidation and corrosion of gas turbine components. Having briefly reviewed the development of gas turbine materials over the past 50 years, the need for corrosion protective coatings and their routes of application and chemistries are explored. The effects of varying coating chemistries and application methods is examined in the context of the major corrosive degradation mechanisms which operate in aircraft and industrial gas turbines. A case study relating to the interdiffusion of coatings and a typical third generation Ni-based alloy is presented which shows that this phenomenon may be of importance with respect to coating life. Finally the paper briefly investigates thermal barrier coatings and how their failure is attributable to the oxidation of the bond coats to which they are attached. Ó 2004 Elsevier Ltd. All rights reserved. Keywords: Gas turbine materials; Aluminide coatings; MCrAlY coatings; Interdiffusion; TBCs
1. Alloy development Since the use of stainless steels for the blades and vanes of the first gas turbine engines, much alloy development has taken place, both with respect to composition and processing [1]. Fig. 1 shows how nickelbased alloy compositions have changed since 1965 and clearly shows that the chromium content has been drastically reduced from about 15 wt% to about 3 wt% and that aluminium contents have increased to about 5 wt%. A key indication from Fig. 1 is the increase in refractory element content, such that over the 30 years period between 1965 and 1995, additions of tantalum, rhenium, tungsten and molybdenum have increased from 8 to 20 wt%. The reasons for these changes in composition are two fold and based on the strengthening mechanisms involved in developing Ni-based alloys for optimum creep resistance at temperatures up to 1100 °C. A typical microstructure of a third generation single crystal Ni-based superalloy is shown in Fig. 2 and clearly shows that the alloy comprises a series of c0 cuboidal precipitates with an interprecipitate ‘‘cement’’ of
*
Fax: +353-61-338172. E-mail address: [email protected] (M.J. Pomeroy).
0261-3069/$ - see front matter Ó 2004 Elsevier Ltd. All rights reserved. doi:10.1016/j.matdes.2004.02.005
a c Ni-matrix. The composition of c0 is effectively Ni3 Al and thus the c matrix contains the refractory alloying elements Ta, Re, W and Mo, although much Ta enters c0 due to its solubility in this phase. The presence of these refractory elements in solid solution in the matrix gives rise to significant solid solution strengthening due to the strength of bonding to Ni [1]. Furthermore, the diffusion effects required for dislocation climb aided creep or diffusion creep are also slowed by the strong interatomic bonding between Ni and the refractory elements [1]. Clearly, as the refractory element content is increased, then alloys become stronger (due to solid solution strengthening) and more creep resistant (due to stronger interatomic bonding). Thus, as shown in Fig. 1 optimised c0 and therefore Al contents were arrived at in the early 1980s and further improvements have only been realised via solid solution strengthening additions. Processing techniques have also led to improved alloy strength and creep resistance with high purity alloys being prepared via vacuum melting and casting techniques. Furthermore, since creep failure is almost always a grain boundary initiated phenomenon due to either grain boundary sliding or vacancy condensation on boundaries perpendicular to the applied tensile stress, the development of directionally solidified or single crystal turbine components has resulted in further creep
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applied to decrease the temperature of the hottest parts of the turbine components by up to 170 °C.
25
wt. % element
20
2. Oxidation and corrosion degradation processes 15
10
5
0 1960 Cr
1970 Ti
1980 year
1990
2000
Al
Mo+Ta+W+Re
Fig. 1. Variations in alloying element additions to Ni-based alloys with year of alloy introduction (based on data in [2]).
In order to understand the development of corrosion protective coatings, it is necessary to appreciate the processes by which oxidation and corrosion occur and how the mechanisms by which they occur depend on environment and temperature. Three accelerated degradative processes occur depending on temperature and these can be defined in order of increasing temperature as: Type II hot corrosion, Type I hot corrosion and oxidation (see Fig. 3). Type II hot corrosion occurs at temperatures in the range 600–850 °C and involves the formation of base metal (nickel or cobalt) sulphates which require a certain partial pressure of sulphur trioxide for their stabilisation. These sulphates react with alkali metal sulphates to form low melting point compounds which prevent a protective oxide forming. Indeed, the oxide formed is a striated one as demonstrated in the work of Viswanathan [3]. As discussed later, Nicholls et al. [4] have shown that Ni–Cr–Al materials resist this form of corrosion most effectively. Type I hot corrosion involves the transport of sulphur from a sulphatic deposit (generally Na2 SO4 ) across a preformed oxide into the metallic material with the formation of the most stable sulphides. Once stable sulphide formers (e.g. Cr) are fully reacted with the sulphur moving across the scale, then base metal sulphides can form with catastrophic consequences as they are molten at the temperatures at which Type I hot corrosion is observed (750–950 °C) [3,5]. Thus, the formation of NiS2 (molten at 645 °C) and Cox Sy (lowest
property improvements. Thus, directionally solidified components are cast such that there are a minimum of grain boundaries perpendicular to the major tensile stress and single crystal components, of course, contain no grain boundaries. All of this alloy development has been made to improve mechanical properties to facilitate higher component temperatures. However, the low chromium and aluminium contents mean that these materials do not have the necessary intrinsic resistance to oxidation and corrosion required for the long term operation associated with gas turbine operation. Accordingly, coatings must be applied to gas turbine components for oxidation or oxidation/corrosion resistance. In addition, due to the demand to increase turbine inlet temperatures and thus cycle efficiencies, ceramic insulating coatings can be
corrosion rate (arb. units)
Fig. 2. Typical microstructure of a Ni-based superalloy showing cuboidal c0 and c matrix (bar ¼ 1 lm).
500
700 900 temperature (˚C)
Type II
Type I
1100
Oxidation
Fig. 3. Schematic representation of rate-temperature curves for Type II hot corrosion, Type I hot corrosion and oxidation (alumina former).
M.J. Pomeroy / Materials and Design 26 (2005) 223–231
log (parabolic rate constant)
liquidus 840 °C) can cause degradation levels which are serious enough to cause major component degradation. The most suitable materials which can resist Type I hot corrosion are PtAl2 –(Ni–Pt–Al) coatings and M–CrAlY coatings containing up to 25 wt% Cr and 6 wt% Al [6]. The oxidation of metals depends on the rates at which anion or cation transport can occur through the crystal lattice or along grain boundaries in the oxide [7]. In alloys, which oxide is the most stable depends on the oxide dissociation pressure which is lowest for Al and Cr compared to the typical base elements Fe, Co and Ni. For Ni–Cr alloys containing more than about 10 at.% Cr, a continuous protective chromia layer is formed. A higher Cr level (25 at.%) is required for Co-based alloys because Cr diffuses more slowly in Co and so cannot form a continuous chromia layer at lower concentrations [7]. Chromium oxide can itself oxidise at temperatures greater than about 850 °C to a volatile CrO3 compound. Because of this, the use of aluminium additions for oxidation resistance is preferred at this temperature and above for key components such as those used in gas turbines. There is added advantage, as shown schematically in Fig. 4, because the rate of formation of aluminium oxide is slower than for chromium oxide at the same temperature. Addition levels of Al to Ni for the formation of a continuous alumina scale are as for chromia formation (10 at.%). Alumina formation at lower Al contents can be induced by mixed Cr + Al additions (see Fig. 4) since Cr ‘‘getters’’ oxygen allowing alumina to form at lower aluminium activities (contents) [7]. The aluminium and chromium contents referred to above apply to isothermal oxidation conditions. However, when thermal cycling conditions prevail, oxide scales can spall from the substrate surface due to thermally induced stresses [8]. The oxidation resistance under such conditions can be markedly improved by the addition of so-called reactive elements (Y, Hf, Ce) to
Cr
Al + Cr Al reciprocal temperature
Fig. 4. Schematic Arrhenius plot for oxidation of alloys containing Cr, Cr + Al and Al.
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alloys and coatings. An exhaustive search for the reasons behind this effect has finally shown that yttrium and rare earth metals segregate to grain boundaries within alumina scales causing a reduction in Al and O transport rates through the oxide and thus reduced oxidation rates [9]. Furthermore, reactive elements combine with sulphur and phosphorus impurities in metallic materials and coatings with the result that these impurities cannot selectively diffuse to the surface and contaminate the oxide–metallic interface. This gives extremely good scale adherence and categorically explains why minor yttrium or cerium (<0.8%) additions to oxidation resistant coatings and alloys greatly improve oxidation resistance.
3. Corrosion resistant coatings-basic processes From the above analysis, aluminium or chromium rich coatings are required for oxidation and corrosion resistance with the former type being required for resistance to oxidation and Type I hot corrosion and the latter type for Type II hot corrosion. Two categories of coatings arise: diffusion coatings and overlay coatings. Diffusion coatings are formed by a chemical vapour deposition type processes e.g. aluminising or chromising or indeed codeposition of both Al and Cr. Whilst these processes are primarily CVD-based, they are more commonly referred to as diffusion coatings since their application involves interdiffusion between deposited Al or Cr and the substrate onto which they are coated. Thus, for example the aluminising of a nickel-based superalloy gives rise to a coating of generic composition NiAl (b-nickel aluminide). Several variants of these ‘‘ising’’ processes arise. These are: pack [10], out of pack [10], CVD [11] and fluidised bed CVD [12]. With pack aluminising or chromising components to be coated are placed in a ‘‘pack’’ of a diluent, typically alumina, a halide activator (e.g. ammonium fluoride) and metal powder e.g. Al or Cr and heat treated in argon. For the out of pack process, the component is located downstream from the ‘‘pack’’ such that the metal halide gas produced by the pack impinges on the substrate and the metal is deposited on the surface where it undergoes interdiffusion with the substrate. In the CVD process Al and HCl are reacted under controlled conditions and passed under a positive pressure into the coating part of the system where they enable aluminide formation [13]. A particular benefit of this system is the cleanliness of the coatings yielding improved oxidation resistance [11] and the fact that the gas can be passed through, as well as over, turbine components with cooling holes such that internal and external surfaces are coated as well as the cooling holes. In contrast, pack aluminising can cause cooling holes to be blocked with diluent particles which require removal after processing. An additional
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benefit of the CVD process is that the chemistry can be carefully controlled enabling additional elements to be incorporated into coatings at precise levels [13]. Such ideas form the basis of an EU R&D Growth project (G4RD-CT-2000-00319). The fluidised bed aluminising process [12] involves the suspension of the component in a bed of inert particles (e.g. alumina) by a gas stream comprising the gases necessary for aluminising in a carrier gas. Given that fluidised bed technology is characterised by excellent solids mixing and heat transfer but poor gas mixing, temperature control may be excellent but inlet gas chemistries must be carefully controlled. Overlay coating involves the application of coatings to substrates using physical deposition processes. Typical application methods include: thermal spraying, physical vapour deposition (PVD) electron beam physical vapour deposition (EBPVD), ion plating/sputtering and electroplating [14].
4. Corrosion resistant coatings-types and chemistries 4.1. Overlay coatings Overlay coatings typically comprise b þ c0 aluminide in a c matrix [1] and are of typical composition (Ni, Co)15–28 wt% Cr, 4–18 wt% Al, 0.5–0.8 wt% Y [6]. The relative amounts of Ni and Co depend upon: (a) coating ductility requirements (>Ni) and (b) corrosion resistance (>Co) [6]. The excellent oxidation and corrosion resistance of the coatings is afforded by the formation of highly adherent alumina scales which grow slowly as a result of the reactive element effect attributable to yttrium. In addition, their high Cr contents make them useful protective coatings against Type II hot corrosion. These overlay coatings are deposited using thermal spray (Ar shrouded plasma (APS) or low pressure plasma (LPPS)) techniques or electron beam physical vapour deposition (EBPVD). The EBPVD technique is typically preferred for high quality coatings since a certain amount of oxidation of the coating particles typically occurs during thermal spraying giving rise to nanoscale oxide particles at splat boundaries. However, the thermal spray processes are frequently used in practice. Overlay coatings show interdiffusion effects with alloy substrates and Itoh and Tamura [15] indicate the rate of interdiffusion decreases in the order NiCrAlY > CoCrAlY > NiCoCrAlY > CoNiAlY, where NiCo represents higher Co contents compared to Ni and CoNi higher Co contents. These results, collected using vacuum plasma sprayed coatings were independent of the three substrates used in the experimental programme. The most recent technological advance in overlay coating technology is the development of so-called smart coatings by Nicholls et al. [4]. These coatings attempt to
address the problems associated with the differences in temperature over the surface of an airfoil. Temperatures vary from a maximum of the order of 1100 °C at leading and trailing edges to about 650 °C in the centre of the airfoil surfaces and near airfoil roots. Because of this, the nature of environmental degradation varies from oxidation through Type I hot corrosion to Type II hot corrosion as indicated in Fig. 2. A practical example of this variation in corrosion type with temperature is graphically portrayed by Viswanathan [5] in his analysis of the failure of airfoils. Following extensive laboratory studies employing alkali–metal sulphates in simulated gaseous environments, Nicholls et al. [4] have been able to very carefully define optimised chromium and aluminium contents for oxidation and corrosion protection in the temperature ranges referred to above. As might be expected, high Cr contents (>40 wt%) and low Al contents (6–8 wt%) are most suitable for protection from Type II hot corrosion (650 °C). For protection from Type I hot corrosion (800–950 °C), roughly equal amounts of Cr and Al represent optimised compositions. Oxidation protection (1100 °C), is best afforded by coatings containing 25 wt% Cr and 14 wt% Al. The final coating developed by Nicholls et al. [4] comprises a commercial base coating (Co–32Ni–21Cr–8Al–0.5Y) adjacent to the substrate, a Cr enriched layer of variable composition from Ni–60Cr–20Al to Ni–35Cr–40Al and a surface layer of composition Ni–15Cr–32Al. These multi-layer coatings have been shown to out perform typical a commercial Pt modified aluminide and an Al enriched version of the base coat at 700–800 °C. Whilst this technology appears a step forward, issues related to coating ductility and additional improvements afforded by cobalt with respect to corrosion resistance, may be of additional importance. 4.2. Diffusion coatings 4.2.1. Chromium rich coatings Effectively, chromising results in Cr enrichment of the surface layers which results in either greater Cr levels in the bulk element solid solution (i.e. Fe, Co, Ni) or supersaturated solid solutions plus aCr phases [16]. As for aluminising, chromising is effected by heating components in a gas rich in a volatile Cr compound (e.g. CrCl3 ). Chromising is of particular use for resistance toward Type II hot corrosion and has also been found to be of major benefit in protecting Ni-based alloys from corrosion by sulphatic deposits in chemical plant. That said, Ni–50Cr alloys have been shown to be equally as corrosion resistant. 4.2.2. Aluminide coatings Aluminising is carried out by two different processes which differ in terms of the activity of aluminium in the gas phase and the temperature at which the aluminising
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process is carried out [10]. The low activity high temperature (LAHT) is a ‘‘one step’’ process as far as the development of a b-NiAl coating is concerned. On the other hand the high activity low temperature route typically gives rise to a d-Ni2 Al3 coating which requires subsequent heat treatment to convert it to b-NiAl. As might be expected, the mechanisms by which the coatings grow is also different. For the LAHT process the aluminium activity is insufficient for it to be the predominant diffusing species and accordingly, coatings form by the diffusion of nickel from the alloy substrate into the region of the coating. For the HALT process the aluminium activity in the gas phase and at the surface of the coating is high enough to facilitate the inward diffusion of Al into the alloy substrate. The differing growth modes may be of importance when considering coating integrity, since coatings which grow by outward Ni diffusion can trap diluent particles (alumina) in aluminising packs within them. Which coating process is selected depends upon a number of features e.g. heat treatment specifications for the substrate alloy, nature of packs available, integrity issues, etc. [10]. 4.2.3. Platinum modified aluminide coatings Platinum modified nickel aluminide coatings exist in two forms, a two phase PtAl2 – (Ni–Pt–Al) (Fig. 5(a)) or single phase Pt modified b-NiAl (Fig. 5(b)). The manner by which the coatings are formed involves an initial deposition of a layer of platinum, typically 6 lm thick either by electrodeposition or ion plating. Following a post coating annealing process to diffusion bond the platinum to the substrate, aluminising gives rise to the platinum modified NiAl coating. If the prealuminising heat treatment is carried out such that significant platinum diffusion into the surface layers of the substrate occurs, the aluminide formed is of the single phase variety. If the pre-aluminising anneal is conducted such that platinum diffusion effects are small, then the bi-phasic PtAl2 – (Ni–Pt–Al) forms [17]. One of the original reasons for platinum modification was to try to suppress the diffusion of refractory elements (W, Mo) into the surface layers of aluminide coatings. Results obtained by Leyens et al. [18] clearly indicate that this effect is not wholly realised and the subsequent work by Zhang et al. [19,20] has shown that the most probable beneficial effect of platinum in aluminide coatings is to suppress void formation at the coating-alumina scale interface. This greatly improves scale adhesion and accordingly oxidation and corrosion resistance. Two clear examples of the improvement of the enhanced oxidation resistance afforded by Pt modification of aluminide coatings are reported by Krishna et al. [17] and Purvis and Warnes [21]. Results from both of these studies show improved resistance to alumina scale spallation as Pt contents are increased.
Fig. 5. Microstructures of Pt modified aluminide coatings on Ni-based superalloy: (a) two phase PtAl2 + (Ni–Pt–Al) coating; (b) single phase (Ni–Pt–Al) coating (bar ¼ 20 lm).
5. Coating-substrate interdiffusion with diffusion aluminide coatings Coatings have a totally different composition from the alloys to which they are applied. For example a stoichiometric NiAl coating contains 31.5 wt% Al and 68.5 wt% Ni. In contrast, as is apparent from Fig. 1, typical Al contents of Ni-based superalloys are 5 wt% and Ni contents range between 60 and 75 wt%. In order to more clearly illustrate this point, Fig. 6 shows concentration profiles for Ni, Pt and Al across the coating and diffusion zones and into the alloy substrate for a two phase Pt modified NiAl coating on a Ni-based superalloy. Because of these concentration differences, coating – substrate interdiffusion effects arise when alloy-coating diffusion couples are put into service. Accordingly, Ni is expected to diffuse from the substrate into the coating and Pt and Al are expected to diffuse countercurrently. Basuki et al. [22] and Chen and Little [23] have studied the effects of such interdiffusion effects on microstructure and observed the degradation of the as coated single phase b-NiAl to b þ c0 (Ni3 Al). In addition, the
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elemental concentration (at.%)
80 70 60 50 40 30 20 10 0 0
50 100 distance from coating surface (µm)
Fig. 6. Compositional variations for Ni(r), Pt(d) and Al(j) across coating, diffusion layer and into alloy for cross-section shown in Fig. 5(a) .
interdiffusion effects caused microstructural change in the substrates. In particular, Chen and Little [23], who used a Pt modified NiAl coating concluded that Pt diffusion into the substrate resulted in the precipitation of topologically close packed (TCP) phases. These phases are known to be deleterious to the mechanical behaviour of Ni-based superalloys and much work has gone into predicting compositional and temperature ranges over which they form such that these compositional and temperature ranges can be avoided [1]. Recent work at the University of Limerick has looked at the interdiffusion of Pt modified aluminide coatings with CMSX-4 and CMSX-10 alloy substrates. Fig. 7 shows typical microstructures of the coatings on the latter substrate after isothermal exposure in air for 750 h at 950 °C and 375 h at 1100 °C. After both exposure periods, the PtAl2 + (Ni–Pt–Al) coating has transformed to a single phase (Ni–Pt–Al) coating and bright grain like features arise in the coating. These light grains have a nickel rich composition and contain tantalum. Since Ta is soluble in c0 (Ni3 Al) but not in b-NiAl, the original Pt modified b-NiAl has transformed to a mixed b þ c0 . This is expected on the basis of the data given in Fig. 8 which presents data for variations in Ni, Pt and Al contents of the exterior layer of the coatings with exposure time. Given that some 36 at.% Al is the minimum required for b to remain stable, it can be seen from this figure that after 375 h at 1100 °C and 750 h at 950 °C the b phase must convert to b þ c0 . There is still however, sufficient Al in the coating to confer oxidation protection. Although this is the case, the Ta rich c0 grains are prone to catastrophic oxidation and thus if the alumina scale formed above them fails then pitting oxidation may arise. Examples of this have been observed after
Fig. 7. Microstructural changes in coating shown in Fig. 4(a) after oxidation exposure at 950 °C for 750 h (a) and 1100 °C for 375 h (b) (bar ¼ 20 lm).
1500 h exposure at 1100 °C. Similar effects have been observed for the CMSX-4 alloy. However, in contrast to the results of Chen and Little [23], this work has shown no evidence of Pt diffusion into the substrate. The results obtained do however complement the results presented by Chen and Little in that Pt concentration gradients rapidly homogenise as suggested by Fig. 8. Close scrutiny of Fig. 7 shows the occurrence of needle like phases in the substrate alloy. These phases are Re, W and Cr rich and X-ray diffraction shows them to be the TCP phase referred to as r-phase. Fig. 9 plots the depth of penetration of these needle like precipitates into the alloy with respect to the square root of exposure time. Two factors are apparent: (a) an incubation period exists at 950 °C such that penetration after 188 h is only slightly greater than in as coated specimens and (b) straight line relationships arise which are indicative of a diffusion controlled microstructural change mechanism. Detailed examination of the alloy microstructure at the root of a needle-like precipitate, shows that the needle grows along a pre-existing c matrix channel. In addition,
concentration (atomic percent)
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78 65 52 39 26 13 0 0
200 400 600 time of isothermal exposure (hours) Pt 950˚C
Pt 1100˚C
Al 950˚C
Al 1100˚C
Ni 950˚C
Ni 1100˚C
800
Fig. 8. Compositional changes in exterior 16 lm of coating after oxidation exposure at 950 and 1100 °C.
depth of penetration of needle phases into substrate (µm)
250
200
150
100
50
0 0
500 1000 1500 2000 square root of isothermal exposure time (s1/2 ) 950˚C
1100˚C
Fig. 9. Effect of exposure time on depth of penetration of needle like phases into CMSX-10 alloy at 950 and 1100 °C.
in contrast to an original alloy microstructure comprising cuboidal c0 precipitates of highly regular shape and size (see Fig. 2), cuboids adjacent to the needle like precipitates appear to have combined to form an envelope around the needle. Away from this envelope, other cuboids have coalesced to form oblong shaped grains a microstructural feature referred to as rafting [24]. Sims et al. [1] draw attention to decreases in creep rupture and impact strength properties associated with TCP phase formation and Feng et al. [25] have shown that rafted microstructures compromise creep rupture properties. Accordingly, the coating-substrate interdiffusion effects shown for the CMSX-10 alloy which are also observed for the CMSX-4 alloy may give rise to a weakening of the substrate superalloy with increasing exposure time (see Fig. 10).
Fig. 10. Detail of development of needle like phase and microstructural rafting in CMSX-10 alloy (bar ¼ 5 lm).
Additional detailed concentration profile determinations after different times of isothermal exposure at 950 and 1100 °C for both CMSX alloys clearly show that Ni diffuses from the substrate into the coating and Al diffuses from the coating into the substrate and that the depth of the diffusion effects into the alloy increases with time. This data, in addition to the microstructure adjacent to the needle like phases, leads to the conclusion that Al diffuses into the bulk of the alloy via c matrix channels and Ni diffuses outward along the same route. Both of these effects lower the Ni activity in the c matrix channels with the result that the activities of solid solution strengthening elements (Re, W, Cr) increase to levels where they become insoluble as explained by Blavette et al. [26]. This insolubility leads to their precipitation at the route of growing r-needles. In areas away from the growing needles, decreases in Ni activity and increases in Al activity within c matrix channels probably approach the c0 (Ni3 Al) stoichiometry. If this occurs then the c channel will be converted into c0 giving rise to what appears to be cuboid coalescence and hence rafting. These interdiffusional effects are clearly long term factors and accordingly may influence the long term stability of both coating and substrate with respect to the functions they are expected to perform. The possibility of decreases in oxidation resistance of the coatings and decreases in mechanical properties of the alloy component may give some cause for concern. That said, there is little evidence of failures due to such interdiffusional effects although long term experience has only been gained with second generation turbine materials. The observations reported above may, however, indicate that further development of aluminide coatings is needed for third generation gas turbine alloys such as CMSX-10 if interdiffusional effects are not to be life limiting.
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6. Thermal barrier coatings The application of 0.1–0.2 mm thick thermally sprayed or EBPVD coatings of partially stabilised zirconia result in decreases in metal surface temperatures of some 170 °C or, for the same metal surface temperatures, increases in turbine inlet temperatures of 170 °C [27]. Brindley and Miller [27] indicate that such increases in turbine inlet temperatures can result in increases in thrust of 5% and efficiency increases of the order of 1% with attendant fuel economies. Partial stabilisation of the zirconia is achieved by additions of yttria and 6–8 wt% levels appear optimum. Whilst thermally sprayed coatings are adequate for stationary components (e.g. nozzle guide vanes), EBPVD coatings are preferred for rotating components. Examples of the structures of these coatings can be found in [28]. The columnar structure of the EBPVD coating facilitates a certain degree of strain tolerance and this is why this coating type are preferred on rotating parts. Thermal barrier coatings (TBCs) are applied to bond coats which were initially simple NiAl coatings. MCrAlY or Pt modified NiAl coatings are now the norm. Failure processes associated with TBCs are invariably associated with their decohesion from the alumina scale which grows on the bond coat. Originally, Al depletion in the bond coatings gave rise to the formation of spinels which due to their faster growth rate and voluminous oxide caused TBCs to spall [29]. A recent paper [30] also draws attention to this failure mode. It is now however becoming apparent that TBC failure is associated with the growth and adhesion properties of the alumina scales formed on MCrAlY and Pt modified NiAl bond coats [31–33]. Clarke et al. [34] and Tolpygo et al. [35] have shown how stresses develop during the growth of alumina and the fact that the surface of alumina scales are rumpled due to enhanced grain boundary growth rates. The stress build ups as well as oxide rumpling cause alumina to decohere from the bond coat as faster growing grain boundary regions of the oxide effectively push the TBC away from flatter regions of the oxide. Due to crack initiation at the TBC-bond coat interface, TBC decohesion occurs. It would seem that the work completed by Kim et al. [36] shows that the type of thermal cycling testing, and therefore service condition, is important with respect to TBC life. This work shows that Pt modified bond coats behave in a superior manner to MCrAlY coats if cycle times are long (>1 h) whereas the opposite is true for short cycle times (10 min).
7. Conclusion From the above discussion of coatings for gas turbines, it is clear that many types arise and which coating is best depends upon the operating conditions in which
it is to be employed. For Types II hot corrosion MCrAlY coatings present the best protection method whilst for Type I hot corrosion Pt rich coatings and MCrAlY coatings appear best. Generally, for oxidation protection and TBC bond coats Pt modified NiAl would appear best particularly if the aluminising is performed using gas CVD aluminising. It does however become quite clear that alloy-coating interdiffusion reactions have received little attention in the scientific literature, thus coatings are typically only tested with respect to their corrosion and oxidation resistance. Accordingly, whilst the current coating-turbine alloy systems are adequate, the drive for aero-turbine CO2 emission reduction by the EU and the improved economics of ultrahigh efficiency combined cycle electricity generation will see the introduction of alloys with high refractory element contents to meet the exacting requirements. It may be that, for these alloys, coating substrate interactions may control life expectancy and accordingly this area of research will become of significant importance.
Acknowledgements The author wishes to thank Mr Michael Reid who is conducting his PhD programme on Ni-based alloyaluminide coating interdiffusion effects for his help in preparing this manuscript. The coating work at the University of Limerick is sponsored by the Materials and Surface Science Institute under the auspices of the Higher Education Authority’s Programme for Research in Third Level Institutions 1999–2002. Thanks are also due to SIFCO Turbine Components Ltd., Carrigtwohill, Co. Cork for their assistance with Mr Reid’s project.
References [1] Sims CT, Stoloff NS, Hagel WC. Superalloys II. New York: Wiley; 1987. [2] Caron P, Khan T. Evolution of Ni-based superalloys for single crystal gas turbine blade applications. Aerospace Sci Technol 1999;3(8):513–23. [3] Viswanathan R. An investigation of blade failures in combustion turbines. Eng Fail Anal 2001;8(5):493–511. [4] Nicholls JR, Simms NJ, Chan W, Evans HE. Smart overlay coatings-concept and practice. Surf Coat Technol 2002;149(23):236–44. [5] Eliaz N, Shemesh G, Latanision RM. Hot corrosion in gas turbine components. Eng Fail Anal 2002;9(1):31–43. [6] Goward GW. Progress in coatings for gas turbine airfoils. Surf Coat Technol 1998;109(1–3):73–9. [7] Wood GC, Stott FH. Oxidation of alloys. Mater Sci Technol 1987;3(7):519–30. [8] Mevrel R. Cyclic oxidation of high-temperature alloys. Mater Sci Technol 1987;3(7):531–5. [9] Pint BA. Experimental observations in support of the dynamic – segregation theory to explain the reactive–element effect. Oxid Met 1996;45(1–2):1–37.
M.J. Pomeroy / Materials and Design 26 (2005) 223–231 [10] Pichoir R, Hauser JM. Current status of coatings for high temperature applications. In: Environmental degradation of high temperature materials. Inst. Metall. Spring Residential Conf., Series 3, Number 13, Vol. 1; 1980. p. 6/1–6/21. [11] Warnes BM, Punola DC. Clean diffusion coatings by chemical vapor deposition. Surf Coat Technol 1997;94-95(1-3):1–6. [12] Voudouris N, Christoglou C, Angelopoulos GN. Formation of aluminide coatings on nickel by a fluidised bed CVD process. Surf Coat Technol 2001;141(2-3):275–82. [13] Squillace A, Bonetti R, Archer NJ, Yeatman JA. The control of the composition and structure of aluminide layers formed by vapour aluminising. Surf Coat Technol 1999;120–121:118–23. [14] Stringer J. Role of coatings in energy producing systems: an overview. Mater Sci Eng 1987;87:1–10. [15] Itoh Y, Tamura M. Reaction diffusion behaviors for interface between Ni-based super alloys and vacuum plasma sprayed MCrAlY coatings. J Eng Gas Turb Power 1999;121(3):476–83. [16] Castle AR, Gabe DR. Chromium diffusion coatings. Int Mater Rev 1999;44(2):37–58. [17] Krishna GR, Das DK, Singh V, Joshi SV. Role of Pt content in the microstructural development and oxidation performance of Pt-aluminide coatings produced using a high-activity aluminizing process. Mater Sci Eng A 1998;251(1–2):40–7. [18] Leyens C, Pint BA, Wright IG. Effect of composition on the oxidation and hot corrosion resistance of NiAl doped with precious metals. Surf Coat Technol 2000;133:15–22. [19] Zhang Y, Haynes JA, Lee WY, Wright IG, Pint BA, Cooley KM, et al. Synthesis and cyclic oxidation behavior of a (Ni, Pt) Al coating on a desulfurized Ni-base superalloy. Metall Mater Trans A 1999;30(10):2679–87. [20] Zhang Y, Haynes JA, Lee WY, Wright IG, Pint BA, Cooley KM, et al. Effects of Pt incorporation on the isothermal oxidation behavior of chemical vapor deposition aluminide coatings. Metall Mater Trans A 2001;32(7):1727–41. [21] Purvis AL, Warnes BM. The effects of platinum concentration on the oxidation resistance of superalloys coated with single-phase platinum aluminide. Surf Coat Technol 2001;146:1–6. [22] Basuki E, Crosky A, Gleeson B. Interdiffusion behaviour in aluminide-coated Rene 80H at 1150 °C. Mater Sci Eng A 1997;224(1-2):27–32. [23] Chen JH, Little JA. Degradation of the platinum aluminide coating on CMSX4 at 1100 °C. Surf Coat Technol 1997;92(1– 2):69–77.
231
[24] Bressers J, Arrell DJ, Ostolaza K, Valles JL. Effect of an aluminide coating on precipitate rafting in superalloys. Mater Sci Eng A 1996;220(1-2):147–54. [25] Feng H, Biermann H, Mughrabi H. Microstructure-based 3D finite element modelling of lattice misfit and long-range internal stresses in creep-deformed nickel-base superalloy single crystals. Mater Sci Eng A 1996;214:1–16. [26] Blavette D, Cadel E, Deconihout B. The role of the atom probe in the study of nickel-based superalloys. Mater Charact 2000;44:133– 7. [27] Brindley WJ, Miller RA. TBCs for better engine efficiency. Adv Mater Process 1989;8:29–36. [28] Miller RA. Thermal barrier coatings for aicraft engines – history and directions, in: Proceedings of the Thermal Barrier Coating Workshop. NASA as NASA Conference Publication 3312; 1935. p. 17–34. [29] Sohn YH, Biederman RR, Sisson Jr RD. Microstructural development in physical vapour-deposited partially stabilized zirconia thermal barrier coatings. Thin Solid Films 1994;250(12):1–7. [30] Shillington EAG, Clarke DR. Spalling failure of a thermal barrier coating associated with aluminum depletion in the bond-coat. Acta Mater 1999;47(4):1297–305, MAR 10 1999. [31] Tolpygo VK, Clarke DR, Murphy KS. Oxidation-induced failure of EB-PVD thermal barrier coatings. Surf Coat Technol 2001; 146:124–31. [32] Evans AG, Mumm DR, Hutchinson JW, Meier GH, Pettit FS. Mechanisms controlling the durability of thermal barrier coatings. Prog Mater Sci 2001;46(5):505–53. [33] Yanar NM, Meier GH, Pettit FS. The influence of platinum on the failure of EBPVD YSZ TBCs on NiCoCrAlY bond coats. Scr Mater 2002;46(4):325–30. [34] Clarke DR, Christensen RJ, Tolpygo V. The evolution of oxidation stresses in zirconia thermal barrier coated superalloy leading to spalling failure. Surf Coat Technol 1997;94-95(1-3):89– 93. [35] Tolpygo VK, Clarke DR. Surface rumpling of a (Ni, Pt) Al bond coat induced by cyclic oxidation. Acta Mater 2000;48(13):3283– 93. [36] Kim GM, Yanar NM, Hewitt EN, Pettit FS, Meier GH. The effect of the type of thermal exposure on the durability of thermal barrier coatings. Scr Mater 2002; 46(7):489–95.