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The mechanical properties of coatings and coated systems
Materials Science and Engineering, A 121 (1989) 633 643
633
The Mechanical Properties of Coatings and Coated Systems* M. 1. WOOD
Engineering Materials Division, ERA Technology, Cleeve Road, Leatherhead, Surrey KT22 7SA (U.K.) (Received February 23, 1989)
Abstract
The increasing mechanical demands being placed on high-temperature gas turbine components has generated interest in establishing a sound understanding of the mechanical behaviour of coatings and coated components. This review covers their bulk and in situ properties and behaviour, primarily under cyclic conditions. The close inter-relationship between bulk coating properties, the presence and nature of structural defects within the coating and the interaction between coating and substrate are emphasised. The role of deformation modelling and fracture mechanics in advancing our understanding and design capabilities is highlighted. 1. Introduction
Coatings have long been used to protect turbine blades from the effects of excessive oxidation and corrosion. An impressive wealth of information is now available for this subject, ranging from basic studies of electrochemisty to endurance field trials in operational service. In comparison, the amount of work which has been undertaken even to determine the mechanical properties of coated superaUoys has been small (see ref. 1 for a recent review). Actual studies of the mechanical properties of the coatings themselves are even scarcer. This in part reflects the lack of any generally perceived need requiring an in-depth study. Various semiquantitative approaches were thought to provide an adequate rationale of the observed behaviour: a wellknown example is the use of the coating's ductile-to-brittle transition temperature (DBTT) curve to explain its resistance to cracking under transient thermal strains--thermomechanical fatigue (the DBTT curve is a plot of the coating's tensile ductility as a function of temperature, for example, Fig. 1). Indeed, various coating devel*Keynote lecture. 0921-5093/89/$3.50
opments have pursued the objective of improving the coating's cyclic life in the engine by modifying its ductility characteristics [2]. However, there is now more of a general realisation that characterising a coating's mechanical response through just one property (ductility) is not adequate. For example, "the relative cracking tendency of two coatings systems has been found to completely reverse on two blades of different design and type of operation" [2]. It is important to know the tensile, creep and fatigue properties as well as the behaviour of "defects" in the coating [3, 4]. Thermomechanical fatigue (TMF) problems have inevitably become greater due to increases in the performance of gas turbines. Whilst it is military aircraft that can be expected to have the most severe operating conditions, the coatings used in industrial and marine gas turbines are not guaranteed to be free from cracking problems [5, 6]. It is against this background of increased interest that this overview is set. Its objectives are to cover the somewhat more quantitative relationships that have been developed between the basic properties of a coating and its mechanical interaction with the substrate, primarily under cyclic conditions.
C o - 2 0 C r ~ 1 2 5 At Y Co 18Cr 30
Co
18Cr
9AI
~1 AI ¥
Y DIFFUSED [ ALUMINIDE
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Fig. 1. DBTT curves for various CoCrA1Y coatings together with a diffused aluminide coating [2]. © Elsevier Sequoia/Printed in The Netherlands
634 TABLE 1 References to coating properties
Property
Reference
Expansion coefficient Young's modulus DBqT" curves Tensile properties Creep properties Stress relaxation Fatigue (a) free-standing (b) insitu
MCrAIY
Aluminide coating
NiAI alloys
4, 9 11 4, 12 See 1 for a review 4, 12, 14-17 4, 15, 19, 20 23
7, 11 7 See 1 for a review
8 13
25 !1,14,24
18, 27 21, 22 11 11,14,24
2. Coating properties In order to go beyond qualitative interpretations of observed behaviour it is necessary to know the mechanical and physical properties of the coatings. A survey of the available references is given in Table 1. Whilst it is not exhaustive, the limited size of the data base is apparent. The only exception is for DBTT curves where a fairly large number of systems have been characterised. Information on aluminide-based coatings is even scarcer than for MCrAIY coatings, primarily because of the experimental difficulties of determining their properties (see, for example, refs. 11 and 26; the physical properties have usually been measured on simulated coatings, e.g. castings of the right composition). Most of the studies of MCrA1Y overlay coatings use test pieces cut from thick (1-3 mm) plasma-sprayed deposits although thick electron-beam evaporated deposits [4] and hot isostatically pressed powder [ 11, 23, 30] compacts have also been used (Table 2). Because of the importance of these properties a brief summary of the salient features is in order: TABLE 2
(i) Expansion coefficients. It is well known that coatings and superalloys have differing expansion coefficients (Fig. 2). In general, MCrAIY coatings have larger expansion coefficients than most superalloys whilst aluminide-based coatings may have values lower than superalloys, but only at the higher temperatures. This property differential manifests itself at low temperatures as a residual stress in the coating, the magnitude of which will also depend on the sequence of processing, heat treatment temperatures and the specific substrate in question. As an example of the size of the residual stresses, a fully processed NiCoCrAIY coating on the single crystal SX60A shown in Fig. 2 is under a tensile residual stress of ca. 200 MPa at 20 °C. For aluminised SX60A the residual stress was ca. 40 MPa, also tensile [10]. (ii) Elastic constants. The Young's modulus of both types of coating is reported to be fairly similar, ca. 170 GPa at room temperature. This is a little lower than a typical polycrystalline superalloy, about 200-220 GPa. However, it should be noted that elastic constants measured on single crystals of NiA1 [13] exhibit two features of importance for aluminide coatings. The first is that for the binary alloy the elastic constants are strong functions of composition. The second point is that NiA1 is elastically anisotropic: e.g. for stoichiometric NiA1 the Young's modulus varies as along (100): 100 GPa along (110): 190 GPa along (111 ): 290 GPa The importance of this anisotropy lies in two physical features of aluminide coatings, one general and one specific to coatings on single crystals. The first is that aluminide coatings often have a
Coating compositions and fabrication methods for property determination
Designation in text
Composition
Fabrication
Reference
PWA 273 PWA 276 PWA 286 CoCrA1Y NiCoCrA1Y NiCoCrAIYTa NiCoCrAIY
Aluminide coating Ni20Col7Crl 3AIY NiCoCrAIY + Hf + Si Co 18Cr8A1Y Ni39Co21Cr8AIY Ni23Co20Cr8.5AIY4Ta Ni23Co 18Cr 12.5AIY
In situ, casting lpps a lpps a lpps a lpps a, igpsb lpps a e b - pvdc
7,11 16,19,25 11,20, 23 10,14,15,24 10,14, 15,17,24,29 12 4
alpps = lOW pressure plasma spray. bigps = inert gas shielded plasma sprayed. Ceb-pvd = electron beam-physical vapour deposition.
635
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×10 6
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g_ 14 0/
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SINGLE
C£YSTAL
12
SX 60A
200
Z,00 600 800 TEMPERATURE (°C)
1000
Fig. 2. Mean thermal expansion coefficients for a superalloy single crystal [10], NiCoCrAIY [10] and an aluminide coating
[7]. grain size which is large relative to the thickness of the coating, hence the coatings will not respond isotropically. The second feature comes from Xray diffraction work on aluminised single crystals from which it appears that the coatings exhibit a strong texture, i.e. preferred orientation [10]. This is perhaps not surprising given the nature of the substrate. However, there does not appear to be any published work on this specific area. If a preferred orientation does exist in components, or test pieces, then any analysis of those items must use the appropriate elastic constants, not isotropic values. (iii) Tensile properties. MCrA1Y coatings have large yield stresses at low temperatures (Fig. 3), comparable with many superalloys. However, above 500-600 °C this decreases rapidly to a very low value. At around the same temperature there is an increase in ductility (cf. DBTT curves) al-
though this seems to occur at a slightly higher temperature (50-100°C) than does the start of the rapid fall in strength [12, 15]. (iv) Creep. The creep strength of MCrA1Y coatings falls off very rapidly with temperature, following the temperature dependence of the yield stress [30]. Creep elongations of up to 400% have been reported at the higher temperatures, deformation being superplastic [12, 16]. However, it must be remembered that in common with all superplastic systems, the rupture strain is a strong function of strain rate, hence stress, at any given temperature (Fig. 4). Measurements of the creep properties of aluminide coatings have been attempted by using very thin substrates, 0.25 mm and 0.13 mm before coating [11]. Whilst some stress relaxation work has been carried out, experimental difficulties have hampered the acquisition of creep data by this route [26]. Despite the current absence of creep data on aluminide coatings, it is instructive to compare MCrAIY coatings and NiA1 data on a strain rate basis (Fig. 5) using a "Larson-Miller"-type approach (after ref. 4: utilising the MonkmanGrant relationship that rupture life is inversely proportional to minimum strain rate). The creep rate disparity between NiA1- and MCrAIY-based alloys is clear. Further to this, the creep strength of NiA1 is increased markedly by alloying, even with small ternary additions ( 1 % - 3 % ) of solute such as molybdenum, tantalum, niobium and hafnium [22]. Indeed, even ternary additions of niobium or hafnium can produce an NiAl-based alloy with a strain rate-stress relationship approaching that of some superalloys. Aluminide coatings on blades often contain 3 % - 5 % C r along with smaller quantities of solute such as molybdenum, hafnium, niobium, tantalum or tungsten from the substrate alloy.
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Fig. 3. Yield stress and ductility of free-standing NiCoCrAIY coating alloy [15].
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STRAIN
RATE
I 10- 2 { s -1 )
Fig. 4. Strain rate dependence of ductility for PWA 276 [16].
636
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20
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28
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Fig. 5. Stress vs. strain rate modified "Larson-Miller" parameter for various coating alloys(with imi. in per cent per hour) [4, 15, 19-22]. (v) Fatigue. Work on free-standing coating alloys is restricted to one study on the overlay PWA 276 at 650 °C and 1050 °C [25]. This temperature range allows a comparison to be made between the coating when it is reasonably strong (490 MPa yield stress) and when it is weak (less than 10 MPa yield stress). The lives at 1050 °C were greater than those of 650 °C on a total strain range basis (Fig. 6). Surface observations also revealed that many cracks became visible half way through the 1050 °C tests: little or no surface damage was observed during the 650 °C tests. The increase in life was attributed primarily to lower crack growth rates at 1050°C where multiple
branching transgranular crack growth was observed. This was unlike the situation at 650 °C when singular, intergranular cracks were found. There appear to be only two reported studies where attention has been paid to identifying the "life to crack initiation" for an in situ coating during a fatigue test [11, 14]. Both these studies were based on coated single-crystal substrates. The broader of these two studies (strain-controlled LCF) again encompassed the two extremes of the coating's behaviour (20 °C strong; 800 °C weak). The cyclic life to crack initiation for plasma-sprayed CoCrAIY and NiCoCrA1Y coatings on [001] oriented single crystals was found to be only weakly temperature dependent (Fig. 7). For aluminide-coated [001] oriented single crystals the life to crack initiation was more markedly reduced upon increasing the temperature from 20°C to 800°C [14,24]. In this study it was possible unambiguously to identify the (surface) site of crack initiation: at both temperatures cracks always formed at unpolished regions of the surface (Fig. 8). (Part of the standard processing cycle for these plasma-sprayed coatings is a mechanical polish. This produces a surface finish which is acceptably smooth aerodynamically, but from which it is not necessary to remove all traces of the original sprayed surface.) The relevance of this observation is that it requires us to decide whether the properties of an in situ coating are being governed entirely by its macroscopic properties or by "defects" and/or processing features and hence crack growth properties. The apparent difference between the fatigue behaviour of the free-standing and in situ coatings with increasing temperature (stronger vs. weaker
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.%
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103 CYCLES
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Fig. 6. Fatigue life of free-standing PWA 276 NiCoCrA1Y [2s].
20°C 800°C
.... +
102
i
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103 104 CYCLES
i
105
Fig. 7. Lifeto crack initiationof in situ MCrAIYcoatings [14].
637
Fig. 8. Crack initiation site during isothermal fatigue test for an in situ NiCoCrAIY coating [10]. respectively) could be due to this observation, or one of several others, e.g. strain rate (affecting ductility and AEin) or the use of a different criterion for failure (a 20 #m surface crack or the failure of a 3 mm diameter test piece).
3. Isothermal fatigue Whilst design engineers are primarily interested in the T M F properties of coated components, it is instructive to consider first the case of isothermal fatigue. This is prudent not only because it is simpler than T M F and enlightening as to the physical processes which are occurring but also because it has been the subject of some misconceptions over the years. There have been a moderate number of studies of the effects of coatings on fatigue life: a wide range of observed behaviour has been found (see ref. 1 for a review). However, most explanations have tended to be rather general, e.g. straining the coating excessively below its DBTT causes it to crack. In part this is because of the absence of coating property data. However, some studies have been more particular in trying to separate coating behaviour from the performance of the coated component, and then relating these observations specifically to coating properties and the imposed strain cycle (see, for example, refs. 11, 14, 24, 25, 28 and 29). As a result of such work it has been possible to draw out three (apparently) separate factors [14] which govern the overall effect of the coating.
(i) The ability of the coating to withstand crack formation and the relationship between this and the coating's structure and mechanical properties. (ii) The behaviour of cracks in the coating once they have reached the coating-substrate interface. Do they arrest or continue propagating? (iii) The rate at which (i) and (ii) can operate relative to crack nucleation and growth processes associated with the substrate. The need to make these distinctions can be seen from the results shown in Fig. 9 (CoCrAIY or NiCoCrA1Y plasma-sprayed [001] oriented single crystals, Table 2). At room temperature (cf. Fig. 9(a)) cracks could be detected in the coating (Fig. 8) after 10%-30% of the life of the coated test piece. The presence of the coating reduced the fatigue properties of the single-crystal substrate. At 800 °C cracks were detected at less than 2% of the life of the coated test piece whilst the coating had little or no effect on the fatigue properties of the single crystal (Fig. 9(b)). This failure of the coating well before the failure of the coated test piece has been observed in other work, e.g. PWA 286 and 273 (Table 2) on the single-crystal PWA 1480 [11]. A small difference was that the life fraction at coating crack initiation at 1038 °C was not quite as low as 2% (Fig. 10). There was no observable difference between the MCrAIY and aluminide coatings in their cracking lives at these high temperatures. A similar plot at 760 °C for PWA 273 aluminised (111) single crystals
638 20°C
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PWA 2 7 3 ALUMINIDE
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0.2
COATED
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0.4 0.6 LIFE FRACTION
i
08
COATING FAILURE i
i
i
102
103
104
(a)
i
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10S
Fig. 10. Longest observed surface crack length vs. life fraction during isothermal LCF [11].
CYCLES
800°C
R= -1
2.0
v =0-SSHz
1.0 0-8
\
06 \ \.
0.4
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(b)
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0.7
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104 CYCLES
103
i
105
Fig. 9. (a) Fatigue life of in situ MCrA1Y coating, the MCrA1Y-coated and the uncoated single crystal SX60A at 20 °C [24]. (b) Fatigue life of in situ MCrA1Y coating, the MCrAlY-coated and the uncoated single crystal SX60A at 800 °C [24].
indicated coating failure as occurring as late as 0.5 life fraction. As previously discussed, the decrease in coating life with increasing temperature may be linked to an increase in the inelastic strain range because of the marked drop in strength. Even simple calculations [24] indicate that the deformation will be primarily elastic at 20 °C and inelastic at 800 °C. This must then outweigh the increase in ductility, large though this can be (it should be recalled that the ductility at high temperatures is a function of strain rate, see Fig. 5). However, such an approach does not address the role of "defects" in initiation cracks. Because of the high surface-tovolume ratio for an in situ coating, the importance of surface events will be greatly increased over
that found for a conventional test piece. Insufficient modelling and experimental work has been done to reconcile the data currently available, although phenomenological approaches do show some promise [ 11, 26]. As the in situ coatings failed more rapidly at the higher temperatures the reason for the diminishing effects produced by the coating on the coated test piece's life arises from a decrease in the crack propagation rate, in particular in the vicinity of the coating-substrate interface. From the work on CoCrAIY-NiCoCrAIY-coated single crystals [14, 24] at 20 °C the few cracks which had formed in the coating had all propagated into the substrate (Fig. 11), causing failure of the test piece. The crack growth planes were also markedly affected by the presence of the coating [14, 30] especially at the lower temperatures. At 800 °C the large number of coating cracks which formed had arrested at the coating-substrate interface (Fig. 12). Test piece failure usually occurred by crack growth from internal features, e.g. porosity. This arrest or retardation at the coating-substrate interface has been observed by others [I1, 31] and found to be dependent upon the type of coating (for a given test condition). For example, cracks in a PWA 273 aluminised single crystal (PWA 1480) propagated from the coating into the substrate during an LCF test at 1038 °C whilst cracks in PWA 286 NiCoCrAIY-coated PWA 1480 arrested at the interface. Since the stress range associated with a given strain range decreases with increasing temperature, the possibility that the lower crack growth rates are linked to the resultant lower AK needs to be excluded. To demonstrate that this is not the case, the number of cycles required to grow a coating crack from initiation to test piece failure is
639
Fig. 1I. Cracks in NiCoCrAIY coating propagating into the substrate during 20 °C LCF test [25].
Fig. 12. Cracks in NiCoCrA1Y coating arrested at interface at 800 °C in LCF test [25].
shown in Fig. 13 as a function of the maximum stress in the substrate. For a crack which has reached the coating-substrate interface the stress intensity value K at the tip of the crack at different temperatures will be directly related to the maxim u m stress in the substrate, everything else being equal. The slower crack growth rates at 800 °C are clear, especially at the lower stresses, indicating that the physical situation at 20 °C and 800 °C is different. As few authors have commented in print on the matter of crack arrest, few explanations have been proffered to resolve the question of propagation vs. arrest. Possible explanations have included retardation of short crack growth through a crack closure effect [30, 32] or increased crack tip blunting at high temperatures reducing the stress intensity factor [30, 33]. A third interpretation [14, 24] concentrates on the way load is partitioned between the coating and the substrate for if the coating cracks, the load it was carrying must be redistributed. Because of constraint factors this
redistribution manifests itself as a stress concentration at the tip of the crack. At low temperatures the coating will be strong and hence any crack will have a large stress intensity at its tip. In contrast, when the coating is weak (high MPa
60O
\
~"'""-
LO0
200
20Oc
--
A[umlnlde
<7
M CrA{V
O
i
10 3
800°C • • i
10k
i
10s
• CYCLES
Fig. 13. Maximum stress in the substrate v s . number of cycles to grow a coating crack to fail the single crystal test piece [14].
640 temperature) it carries little load anyway, hence the presence of a crack requires little load to be redistributed. Thus the stress intensity at the tip of a geometrically equivalent crack will be far smaller than when the coating was strong. The crack growth rate beyond the coating will therefore decrease with increasing temperature when the coating softens significantly. Although less explicitly stated than in ref. 14, a somewhat similar observation has been made [31] that non-propagating cracks behave as they do because they are "not sufficiently severe from a stress intensity standpoint to reduce the l07 uncoated strength of these superalloys". This is still a qualitative argument. Verification requires detailed mathematical modelling of the load repartitioning behaviour using constitutive equations to describe the coating's properties. Whilst this approach has been used to model TMF it does not yet seem to have been applied to the simpler case of isothermal fatigue. However, it can be seen that of the three factors identified above, crack initiation and propagation are linked through their common dependence on the mechanical properties of the coating. The third factor, the relative rate at which substrate and coating progress towards failure, is worth commenting on since it illustrates a difference between polycrystalline and single crystal materials. Several published results using straincontrolled LCF show that coatings have little effect on the lives of polycrystalline superalloys (see, for example, ref. 34). Polycrystalline superalloys have shorter lives than [001] oriented single crystals when fatigued under strain control because of the larger stresses induced by their greater Young's modulus. However, since the strain-stress range experienced by the coating is not dependent, to a first approximation, on the modulus of the substrate, its intrinsic life will be independent of substrate type. Hence it can be anticipated that coatings will affect the lives of polycrystalline superalloys less than those of single crystals or unidirectionally solidified superalloys, as is observed. This illustration emphasises the need to use care when interpreting test results on such "composite" systems.
4. Thermo-mechanical fatigue Simple experimental data on the TMF properties of coated and uncoated superalloys is fairly readily available, e.g. based on tests in hot gas
streams or fluidised beds. However, because of the somewhat empirical nature of the data, detailed analyses have not been forthcoming. Various interpretations have been advanced which offer semiquantitative rationales of the observed behaviour, e.g. the use of the DBTT curve. It is only recently that detailed attempts have been made to analyse and model the deformation and failure of the coating under TMF conditions. Various forms of TMF cycles can be postulated. Those most commonly employed are shown in Fig. 14 together with the effect of thermal expansion mismatch on the strain range for the coating under out-of-phase cycling. However, it is generally agreed that the most damaging cycle is that which imposes a tensile strain on the coating at low temperature (out-of-phase cycle): e.g. the cyclic life for NiCoCrAIY and aluminised PWA 1480 single crystals increased from ca. 500-1500 cycles (respectively) for out-of-phase TMF to ca. 10 500 cycles for in-phase TMF. It is this out-of-phase cycle which is the most relevant test vehicle in attempts to analyse the performance of coatings for gas turbine blading. One approach to the problem has been phenomenological. This treats the coating as being subject to the same descriptive rules that have been found useful in the analysis of structural alloys, e.g. inelastic strain range, strain range partitioning or hysteretic energy. However, because a coated component is a composite system where failure occurs as the consequence of a series of sequential events, the life of a coated component, Nf, has to be expressed as [11] Nf=Nc+N~c+Nsp
where N¢ is the cycles to initiate a crack through the coating, N~ is the cycles for coating crack to IN PHASE
z +
J
OUT OF PHASE
J
< cc
OUT OF PHASE +COATING MISMATCH ctc
BITHERMAL
+ IT -
>cxs\"
222"'-'" %---',_~co.,~G SUa~TR~
Fig. 14. Various forms of TMF cycling: in-phase, out-ofphase, bithermal and out-of-phasewith allowancefor coating expansion mismatch strain.
641
penetrate a small distance into the substrate and N~p is the cycles to propagate the substrate crack to failure. The best descriptive model for Arc was found to be that of tensile hysteretic energy [26]
O 1~
where f~ (back stress) and K (drag stress) are strain history dependent internal state variables describing kinematic and isotropic cyclic hardening respectively. Using this approach [26] for Arc it is possible to correlate isothermal overlay coating lives and coating lives under TMF conditions to within a factor of about 2.5 (Fig. 16). Determination of the N~ and N~p terms in the cycle life equation remains to be addressed [11, 26]. Likewise, the aluminide coating model still needs to be developed [26]. One of the objectives of any such work is to predict lives, but it also ought to be capable of indicating why different coatings behave differently. It was noted in the introduction that a
0600 L00
0
,oo yI DATA
-
C/3
~200 I o,"
/
/
427-1149"C
D [3
Y
/o)/ 102
.
102
,
,
103
10 ~"
ACTUAL LIFE
(CYCLES)
Fig. 16. Overlay coating life model correlation of isothermal life data and prediction of TMF lives [26].
major failing of an approach such as the DBTT route was its inability to explain why the relative TMF resistance of different coatings should be cycle- or blade-dependent [2]. A simplistic interpretation of the model would be that the "more ductile" coating would always be best. This is not found in practice, nor in TMF tests, although the data [37, 38] is as yet somewhat limited (Fig. 17). Here there appears to be a switch over in relative performance with the NiCoCrAIY PWA 286 being best at high strain ranges and the PWA 273 aluminide being (potentially) superior at low strain ranges. As yet no mechanistic reason for this difference appears to have been put forward, partly because of the lack of constitutive models for an aluminide coating. A somewhat different approach to the problem has placed more emphasis on the role of defects within the coating [4, 39] based on work with eb-pvd NiCoCrA1Y coatings on a variety of substrates. Once the extra strain induced in the coating by thermal expansion mismatch had been taken into consideration, coating cracking lives
~
PREDICTION ~OC) / O -
v
,~10 z
g
}--
z
0 -200
D TMF- 427 1038~C
1o
where A Wt is the tensile hysteretic energy, v is a complex time and temperature-dependent frequency term [35] and b and m are constants. The values of A Wt are derived from calculated hysteresis loops for the coating. These loops are, in turn, based on constitutive models for the coating. For PWA 286, NiCoCrA1Y, the best agreement between an experimental hysteresis loop and a constitutive model (Fig. 15) was found for the isotropic unified viscoplastic formulation [11, 36]
LCF
z.27760,927,038°C
N~ = C(A lYt) -bvm
O_
ISOTHERMAL
0k ~
J
OUCTP~F-P H A S E o
o;2 STRAIN
66
(%)
0.1
102
, L '
10 3
0
NICoCrAiY PWA 286
•
A l u m l n l d e PWA 273 , , __
,
10 ~
50"/. LOAD DROP LIFE (CYCLES)
Fig. 15. Comparison between experimental and calculated hysteresis loop (unified viscoplastic) for PWA 286 under TMF cycling [26].
Fig. 17. Total strain range v s . life for coated PWA 1480 single crystal for an out-of-phase TMF cycle: 427-1038 °C [37].
642 06
5. Concluding remarks
~z~ :E Lu I< P(z Dm
0/.
QZ)~O
SUBSTRATE 0 MAR-M200÷Hf
" 0 18 %
B 1900+Hf
- 0 18"/o
[] NX-188
~
o~ 10 2
J
,
t
~E~
~ L I 11103
- 0 29% ,
,
,
5x103
CYCLES TO CRACK INITIATION
Fig. 18. NiCoCrAIY coating cracking life at defects for an out-of-phase TMF cycle (427-1038°C) vs. mechanical + thermal mismatch strain [39].
could be correlated with the total strain range for out-of-phase cycling (Fig. 18). An important point arising from this, apart from the role of the expansion mismatch strain, is that initiation occurred at flaws. Constitutive models are based on "bulk" property data. A similar difficulty was noted in the isothermal fatigue section between "bulk" and in situ properties. This again raises the question of the importance of the surface condition and nature of structural flaws or weakness in such a high surface-to-volume "test piece" as an in situ coating. Whilst plasma-sprayed coatings will not have the same structural features as eb-pvd coatings, they are not metallurgically perfect products. It has been observed in the work on plasma-sprayed NiCoCrA1Y PWA 286 that crack initiation occurs at many sites over the coating, with cracks from many sites contributing to produce the failure. In contrast, failure in PWA 273-aluminised single-crystal PWA 1480 was dominated by long circumferential cracks [37]. The question of coating crack growth into the substrate has also been analysed [40], the coating crack being regarded as a geometrical extension of any (short) substrate crack, its effect being somewhat modified by differences in elastic constants and differential expansion. However, as was the case for the equivalent isothermal case, the coating crack cannot be treated purely as a geometric feature: its load-carrying capacity must be taken into account for this will affect AK at the crack tip. This will depend upon the temperature dependence of the coating's properties as well as the rate at which the TMF cycle is carried out and will greatly complicate any such analysis, especially for "slow" thermal transients. In addition to all these factors, environmental effects must be included, for TMF lives in vacuum are greater than those in air [29].
This brief survey of the mechanical properties of coatings and coated components has been very selective. Many factors have not been discussed, e.g. the effects of oxidation and corrosion, interdiffusion and ageing, or the effects of different coating techniques on the properties of MCrA1Y coatings. These are all additional effects which, whilst of importance, further complicate an already complex subject. Although it is clear that advances in understanding and component lifing are being made, there are many topics which are unresolved or simply unclear because of the lack of data, e.g. the relative merits in a high surfaceto-volume material such as a coating of a phenomenological approach based on "bulk" properties vs. an analysis based on crack initiation at surface flaws. However, any fracture mechanics type of approach is likely to be complex both because of the size of the defect but also because of the huge range of coating properties encompassed during a single TMF cycle. The answers to other important questions, such as the causes of the differences between MCrA1Y and aluminide coatings, is hampered by the lack of suitable data, in this case primarily because of experimental difficulties. Whilst deformation modelling has been actively pursued, the basic metallurgical side of the problem seems to have been somewhat overshadowed, e.g. there has been minimal work on crack initiation and growth. Although it is important to be able to life a particular coating or component, there is also the need to know how to modify it metallurgically in order to develop a better coating or improve deposition techniques. The increased effort and interest shown in this topic over recent years has brought into clearer focus some of the physical and metallurgical factors governing coating life. However, there is a great deal of work needing to be done before a sound quantitative understanding has been achieved.
Acknowledgment The author is grateful to the Directors of ERA Technology Ltd. for permission to publish this article.
References 1 A. Strang and E. Lang, High Temperature Alloys for Gas Turbines, Liege, October 1982, Reide], Dordrecht, 1982, p. 469.
643 2 G. W. Meetham, Mater. Sci. Technol., 2 (1986) 290. 3 A. L. Bement, High Temperature, Theory and Design, Metallurgical Society of AIME, Warrendale, PA, 1984, p. 2. 4 T. E. Strangman, Ph.D. Thesis, University of Connecticut, 1978. 5 M. I. Wood and J. E. Restall, U.K. Corrosion '87, Brighton, October 1987. 6 R. Burgel, Mater. Sci. Technol., 2 (1986) 302. 7 M. Kaufman, TMS Paper F82-11, Metallurgical Society of AIME, Warrendale, PA. 8 R. W. Clark and J. D. Whittenberger, Proc. 8th Int. Thermal Expansion Syrup., 1984, Plenum, New York, p. 189. 9 L. Pawlowski, D. Lombard, F. Tourenne, F. Kassabji and P. Francis, High Temp. High Pressures, 17 (1985) 611. 10 M. I. Wood, RAE TR 88-021, 1988 (Royal Aircraft Establishment, U.K.) 11 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, 2nd Annual Status Rep., NASA-CR-179594, 1987 (National Aeronautics and Space Administration, Washington, DC). 12 J. M. Veys and R. Mevrel, Mater. Sci. Eng., 88(1987) 253. 13 N. Rusovic and H. Warlimont, Phys. Status Solidi A, 44 (1977) 609. 14 M. I. Wood, A S M Europe Conference, Paris, France, September, 1987. 15 M. I. Wood, RAE Tech. Rep. (Royal Aircraft Establishment, Pyestock, Hampshire, U.K.). 16 M. G. Hebsur and R. V. Miner, Mater. Sci. Eng., 83 (1986) 239. 17 T. A. Taylor, M. P. Over, B. J. Gill and R. C. Tucker Jr., J. Vac. Sci. Technol., A3 (1985) 2526. 18 A. G. Rozner and R. J. Wasilewski, J. Inst. Met., 94 (1966) 169. 19 M. G. Hebsur and R. V. Miner, Thin Solid Films, 147 (1987) 143. 20 G. A. Swanson, I. Linask, D. M. Nissley, P. P. Norris, T. G. Meyer and K. P. Walker, 1st Annual Status Rep., NASA -CR- 174952, 1986 (National Aeronautics and Space Administration, Washington, DC). 21 J. D. Whittenberger, J. Mater. Sci., 22 (1987) 394. 22 K. Vedula, V. Pathare, I. Aslandis and R. M. Titran, Mater. Res. Syrup., 39 (985) 411. 23 D. M. Nissley and G. A. Swanson, Nonlinear Constitutive Relationships for High Temperature Applications, Akron,
24 25 26
27 28 29 30 31 32 33 34 35 36 37 38 39 40
OH, 1986, NASA-CP-IO010, National Aeronautics and Space Administration, Washington, DC, p. 419. M. I. Wood, RAE TR 88-057, 1988 (Royal Aircraft Establishment, U.K.). J. Gayda, T. P. Gadd and R. V. Miner, Int. J. Fatigue, 8 (1986) 217. G. R. Halford, T. G. Meyer, R. S. Nelson and D. M. Nissley, 33rd Int. Gas Turbine and Aeroengine Congress, Amsterdam, The Netherlands, June 1988, ASME, New York. E. M. Schulson and D. R. Barker, Scr. Metall., 17(1983) 519. D. E. Graham and D. A. Woodford, Metall. Trans. A, 12 (1980) 329. J. Gayda, T. P. Gabb, R. V. Miner and G. R. Halford, 1987 TMS AIME Annual Meeting, Denver, Colorado, February 1987. M. 1. Wood and J. E. Restall, High Temperature Alloys for Gas Turbines and Other Applications, Liege, October 1986, Reidel, Dordrecht, 1986, p. 1215. T. E. Strangmen, M. Fuji and X. Ngugen-Dinh, 5th Int. Conf. on Supralloys, Seven Springs, U.S.A. 1984, Metallurgical Society of AIME, Warrendale, PA, p. 797. S. Suresh, C. F. Zaminski and R. O. Ritchie, Metall. Trans. A, 12 (1981) 1435. J. S. Crompton, S. A. Morley and J. W. Martin, 5th Int. Conf. on Superalloys, Seven Springs, U.S.A., 1984, Metallurgical Society of AIME, Warrendale, PA, p. 771. A. Strang, Proc. Int. Conf. on the Behaviour of High Temperature Alloys in Aggressive Environments, Petten, October 1979, Metals Society, London, 1979, p. 595. W. J. Ostergren, J. Test. Eval., 4 (1976) 304. K. P. Walker, NASA-CR-165533, 1981 (National Aeronautics and Space Administration, Washington, DC). G. A. Swanson and R. C. Bill, 21st Joint Propulsion Conference, Monterey, CA, July 1985, A I A A - S A E ASME-ASEE, Paper AIAA-85-1421. K. R. Bain, 21st Joint Propulsion Conference, Monterey, CA, July 1985, A I A A - S A E - A S M E - A S E E , Paper AIAA85-1366. G. R. Leverant, T. E. Strangman and B. S. Langer, 3rd Int. Conf. on Superalloys 1976, Seven Springs, U.S.A., Claitors, 1976, p. 285. T. E. Strangman and S. W. Hopkins, Ceram. Bull., 55 (1976) 304.
633
The Mechanical Properties of Coatings and Coated Systems* M. 1. WOOD
Engineering Materials Division, ERA Technology, Cleeve Road, Leatherhead, Surrey KT22 7SA (U.K.) (Received February 23, 1989)
Abstract
The increasing mechanical demands being placed on high-temperature gas turbine components has generated interest in establishing a sound understanding of the mechanical behaviour of coatings and coated components. This review covers their bulk and in situ properties and behaviour, primarily under cyclic conditions. The close inter-relationship between bulk coating properties, the presence and nature of structural defects within the coating and the interaction between coating and substrate are emphasised. The role of deformation modelling and fracture mechanics in advancing our understanding and design capabilities is highlighted. 1. Introduction
Coatings have long been used to protect turbine blades from the effects of excessive oxidation and corrosion. An impressive wealth of information is now available for this subject, ranging from basic studies of electrochemisty to endurance field trials in operational service. In comparison, the amount of work which has been undertaken even to determine the mechanical properties of coated superaUoys has been small (see ref. 1 for a recent review). Actual studies of the mechanical properties of the coatings themselves are even scarcer. This in part reflects the lack of any generally perceived need requiring an in-depth study. Various semiquantitative approaches were thought to provide an adequate rationale of the observed behaviour: a wellknown example is the use of the coating's ductile-to-brittle transition temperature (DBTT) curve to explain its resistance to cracking under transient thermal strains--thermomechanical fatigue (the DBTT curve is a plot of the coating's tensile ductility as a function of temperature, for example, Fig. 1). Indeed, various coating devel*Keynote lecture. 0921-5093/89/$3.50
opments have pursued the objective of improving the coating's cyclic life in the engine by modifying its ductility characteristics [2]. However, there is now more of a general realisation that characterising a coating's mechanical response through just one property (ductility) is not adequate. For example, "the relative cracking tendency of two coatings systems has been found to completely reverse on two blades of different design and type of operation" [2]. It is important to know the tensile, creep and fatigue properties as well as the behaviour of "defects" in the coating [3, 4]. Thermomechanical fatigue (TMF) problems have inevitably become greater due to increases in the performance of gas turbines. Whilst it is military aircraft that can be expected to have the most severe operating conditions, the coatings used in industrial and marine gas turbines are not guaranteed to be free from cracking problems [5, 6]. It is against this background of increased interest that this overview is set. Its objectives are to cover the somewhat more quantitative relationships that have been developed between the basic properties of a coating and its mechanical interaction with the substrate, primarily under cyclic conditions.
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634 TABLE 1 References to coating properties
Property
Reference
Expansion coefficient Young's modulus DBqT" curves Tensile properties Creep properties Stress relaxation Fatigue (a) free-standing (b) insitu
MCrAIY
Aluminide coating
NiAI alloys
4, 9 11 4, 12 See 1 for a review 4, 12, 14-17 4, 15, 19, 20 23
7, 11 7 See 1 for a review
8 13
25 !1,14,24
18, 27 21, 22 11 11,14,24
2. Coating properties In order to go beyond qualitative interpretations of observed behaviour it is necessary to know the mechanical and physical properties of the coatings. A survey of the available references is given in Table 1. Whilst it is not exhaustive, the limited size of the data base is apparent. The only exception is for DBTT curves where a fairly large number of systems have been characterised. Information on aluminide-based coatings is even scarcer than for MCrAIY coatings, primarily because of the experimental difficulties of determining their properties (see, for example, refs. 11 and 26; the physical properties have usually been measured on simulated coatings, e.g. castings of the right composition). Most of the studies of MCrA1Y overlay coatings use test pieces cut from thick (1-3 mm) plasma-sprayed deposits although thick electron-beam evaporated deposits [4] and hot isostatically pressed powder [ 11, 23, 30] compacts have also been used (Table 2). Because of the importance of these properties a brief summary of the salient features is in order: TABLE 2
(i) Expansion coefficients. It is well known that coatings and superalloys have differing expansion coefficients (Fig. 2). In general, MCrAIY coatings have larger expansion coefficients than most superalloys whilst aluminide-based coatings may have values lower than superalloys, but only at the higher temperatures. This property differential manifests itself at low temperatures as a residual stress in the coating, the magnitude of which will also depend on the sequence of processing, heat treatment temperatures and the specific substrate in question. As an example of the size of the residual stresses, a fully processed NiCoCrAIY coating on the single crystal SX60A shown in Fig. 2 is under a tensile residual stress of ca. 200 MPa at 20 °C. For aluminised SX60A the residual stress was ca. 40 MPa, also tensile [10]. (ii) Elastic constants. The Young's modulus of both types of coating is reported to be fairly similar, ca. 170 GPa at room temperature. This is a little lower than a typical polycrystalline superalloy, about 200-220 GPa. However, it should be noted that elastic constants measured on single crystals of NiA1 [13] exhibit two features of importance for aluminide coatings. The first is that for the binary alloy the elastic constants are strong functions of composition. The second point is that NiA1 is elastically anisotropic: e.g. for stoichiometric NiA1 the Young's modulus varies as along (100): 100 GPa along (110): 190 GPa along (111 ): 290 GPa The importance of this anisotropy lies in two physical features of aluminide coatings, one general and one specific to coatings on single crystals. The first is that aluminide coatings often have a
Coating compositions and fabrication methods for property determination
Designation in text
Composition
Fabrication
Reference
PWA 273 PWA 276 PWA 286 CoCrA1Y NiCoCrA1Y NiCoCrAIYTa NiCoCrAIY
Aluminide coating Ni20Col7Crl 3AIY NiCoCrAIY + Hf + Si Co 18Cr8A1Y Ni39Co21Cr8AIY Ni23Co20Cr8.5AIY4Ta Ni23Co 18Cr 12.5AIY
In situ, casting lpps a lpps a lpps a lpps a, igpsb lpps a e b - pvdc
7,11 16,19,25 11,20, 23 10,14,15,24 10,14, 15,17,24,29 12 4
alpps = lOW pressure plasma spray. bigps = inert gas shielded plasma sprayed. Ceb-pvd = electron beam-physical vapour deposition.
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[7]. grain size which is large relative to the thickness of the coating, hence the coatings will not respond isotropically. The second feature comes from Xray diffraction work on aluminised single crystals from which it appears that the coatings exhibit a strong texture, i.e. preferred orientation [10]. This is perhaps not surprising given the nature of the substrate. However, there does not appear to be any published work on this specific area. If a preferred orientation does exist in components, or test pieces, then any analysis of those items must use the appropriate elastic constants, not isotropic values. (iii) Tensile properties. MCrA1Y coatings have large yield stresses at low temperatures (Fig. 3), comparable with many superalloys. However, above 500-600 °C this decreases rapidly to a very low value. At around the same temperature there is an increase in ductility (cf. DBTT curves) al-
though this seems to occur at a slightly higher temperature (50-100°C) than does the start of the rapid fall in strength [12, 15]. (iv) Creep. The creep strength of MCrA1Y coatings falls off very rapidly with temperature, following the temperature dependence of the yield stress [30]. Creep elongations of up to 400% have been reported at the higher temperatures, deformation being superplastic [12, 16]. However, it must be remembered that in common with all superplastic systems, the rupture strain is a strong function of strain rate, hence stress, at any given temperature (Fig. 4). Measurements of the creep properties of aluminide coatings have been attempted by using very thin substrates, 0.25 mm and 0.13 mm before coating [11]. Whilst some stress relaxation work has been carried out, experimental difficulties have hampered the acquisition of creep data by this route [26]. Despite the current absence of creep data on aluminide coatings, it is instructive to compare MCrAIY coatings and NiA1 data on a strain rate basis (Fig. 5) using a "Larson-Miller"-type approach (after ref. 4: utilising the MonkmanGrant relationship that rupture life is inversely proportional to minimum strain rate). The creep rate disparity between NiA1- and MCrAIY-based alloys is clear. Further to this, the creep strength of NiA1 is increased markedly by alloying, even with small ternary additions ( 1 % - 3 % ) of solute such as molybdenum, tantalum, niobium and hafnium [22]. Indeed, even ternary additions of niobium or hafnium can produce an NiAl-based alloy with a strain rate-stress relationship approaching that of some superalloys. Aluminide coatings on blades often contain 3 % - 5 % C r along with smaller quantities of solute such as molybdenum, hafnium, niobium, tantalum or tungsten from the substrate alloy.
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Fig. 5. Stress vs. strain rate modified "Larson-Miller" parameter for various coating alloys(with imi. in per cent per hour) [4, 15, 19-22]. (v) Fatigue. Work on free-standing coating alloys is restricted to one study on the overlay PWA 276 at 650 °C and 1050 °C [25]. This temperature range allows a comparison to be made between the coating when it is reasonably strong (490 MPa yield stress) and when it is weak (less than 10 MPa yield stress). The lives at 1050 °C were greater than those of 650 °C on a total strain range basis (Fig. 6). Surface observations also revealed that many cracks became visible half way through the 1050 °C tests: little or no surface damage was observed during the 650 °C tests. The increase in life was attributed primarily to lower crack growth rates at 1050°C where multiple
branching transgranular crack growth was observed. This was unlike the situation at 650 °C when singular, intergranular cracks were found. There appear to be only two reported studies where attention has been paid to identifying the "life to crack initiation" for an in situ coating during a fatigue test [11, 14]. Both these studies were based on coated single-crystal substrates. The broader of these two studies (strain-controlled LCF) again encompassed the two extremes of the coating's behaviour (20 °C strong; 800 °C weak). The cyclic life to crack initiation for plasma-sprayed CoCrAIY and NiCoCrA1Y coatings on [001] oriented single crystals was found to be only weakly temperature dependent (Fig. 7). For aluminide-coated [001] oriented single crystals the life to crack initiation was more markedly reduced upon increasing the temperature from 20°C to 800°C [14,24]. In this study it was possible unambiguously to identify the (surface) site of crack initiation: at both temperatures cracks always formed at unpolished regions of the surface (Fig. 8). (Part of the standard processing cycle for these plasma-sprayed coatings is a mechanical polish. This produces a surface finish which is acceptably smooth aerodynamically, but from which it is not necessary to remove all traces of the original sprayed surface.) The relevance of this observation is that it requires us to decide whether the properties of an in situ coating are being governed entirely by its macroscopic properties or by "defects" and/or processing features and hence crack growth properties. The apparent difference between the fatigue behaviour of the free-standing and in situ coatings with increasing temperature (stronger vs. weaker
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Fig. 7. Lifeto crack initiationof in situ MCrAIYcoatings [14].
637
Fig. 8. Crack initiation site during isothermal fatigue test for an in situ NiCoCrAIY coating [10]. respectively) could be due to this observation, or one of several others, e.g. strain rate (affecting ductility and AEin) or the use of a different criterion for failure (a 20 #m surface crack or the failure of a 3 mm diameter test piece).
3. Isothermal fatigue Whilst design engineers are primarily interested in the T M F properties of coated components, it is instructive to consider first the case of isothermal fatigue. This is prudent not only because it is simpler than T M F and enlightening as to the physical processes which are occurring but also because it has been the subject of some misconceptions over the years. There have been a moderate number of studies of the effects of coatings on fatigue life: a wide range of observed behaviour has been found (see ref. 1 for a review). However, most explanations have tended to be rather general, e.g. straining the coating excessively below its DBTT causes it to crack. In part this is because of the absence of coating property data. However, some studies have been more particular in trying to separate coating behaviour from the performance of the coated component, and then relating these observations specifically to coating properties and the imposed strain cycle (see, for example, refs. 11, 14, 24, 25, 28 and 29). As a result of such work it has been possible to draw out three (apparently) separate factors [14] which govern the overall effect of the coating.
(i) The ability of the coating to withstand crack formation and the relationship between this and the coating's structure and mechanical properties. (ii) The behaviour of cracks in the coating once they have reached the coating-substrate interface. Do they arrest or continue propagating? (iii) The rate at which (i) and (ii) can operate relative to crack nucleation and growth processes associated with the substrate. The need to make these distinctions can be seen from the results shown in Fig. 9 (CoCrAIY or NiCoCrA1Y plasma-sprayed [001] oriented single crystals, Table 2). At room temperature (cf. Fig. 9(a)) cracks could be detected in the coating (Fig. 8) after 10%-30% of the life of the coated test piece. The presence of the coating reduced the fatigue properties of the single-crystal substrate. At 800 °C cracks were detected at less than 2% of the life of the coated test piece whilst the coating had little or no effect on the fatigue properties of the single crystal (Fig. 9(b)). This failure of the coating well before the failure of the coated test piece has been observed in other work, e.g. PWA 286 and 273 (Table 2) on the single-crystal PWA 1480 [11]. A small difference was that the life fraction at coating crack initiation at 1038 °C was not quite as low as 2% (Fig. 10). There was no observable difference between the MCrAIY and aluminide coatings in their cracking lives at these high temperatures. A similar plot at 760 °C for PWA 273 aluminised (111) single crystals
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indicated coating failure as occurring as late as 0.5 life fraction. As previously discussed, the decrease in coating life with increasing temperature may be linked to an increase in the inelastic strain range because of the marked drop in strength. Even simple calculations [24] indicate that the deformation will be primarily elastic at 20 °C and inelastic at 800 °C. This must then outweigh the increase in ductility, large though this can be (it should be recalled that the ductility at high temperatures is a function of strain rate, see Fig. 5). However, such an approach does not address the role of "defects" in initiation cracks. Because of the high surface-tovolume ratio for an in situ coating, the importance of surface events will be greatly increased over
that found for a conventional test piece. Insufficient modelling and experimental work has been done to reconcile the data currently available, although phenomenological approaches do show some promise [ 11, 26]. As the in situ coatings failed more rapidly at the higher temperatures the reason for the diminishing effects produced by the coating on the coated test piece's life arises from a decrease in the crack propagation rate, in particular in the vicinity of the coating-substrate interface. From the work on CoCrAIY-NiCoCrAIY-coated single crystals [14, 24] at 20 °C the few cracks which had formed in the coating had all propagated into the substrate (Fig. 11), causing failure of the test piece. The crack growth planes were also markedly affected by the presence of the coating [14, 30] especially at the lower temperatures. At 800 °C the large number of coating cracks which formed had arrested at the coating-substrate interface (Fig. 12). Test piece failure usually occurred by crack growth from internal features, e.g. porosity. This arrest or retardation at the coating-substrate interface has been observed by others [I1, 31] and found to be dependent upon the type of coating (for a given test condition). For example, cracks in a PWA 273 aluminised single crystal (PWA 1480) propagated from the coating into the substrate during an LCF test at 1038 °C whilst cracks in PWA 286 NiCoCrAIY-coated PWA 1480 arrested at the interface. Since the stress range associated with a given strain range decreases with increasing temperature, the possibility that the lower crack growth rates are linked to the resultant lower AK needs to be excluded. To demonstrate that this is not the case, the number of cycles required to grow a coating crack from initiation to test piece failure is
639
Fig. 1I. Cracks in NiCoCrAIY coating propagating into the substrate during 20 °C LCF test [25].
Fig. 12. Cracks in NiCoCrA1Y coating arrested at interface at 800 °C in LCF test [25].
shown in Fig. 13 as a function of the maximum stress in the substrate. For a crack which has reached the coating-substrate interface the stress intensity value K at the tip of the crack at different temperatures will be directly related to the maxim u m stress in the substrate, everything else being equal. The slower crack growth rates at 800 °C are clear, especially at the lower stresses, indicating that the physical situation at 20 °C and 800 °C is different. As few authors have commented in print on the matter of crack arrest, few explanations have been proffered to resolve the question of propagation vs. arrest. Possible explanations have included retardation of short crack growth through a crack closure effect [30, 32] or increased crack tip blunting at high temperatures reducing the stress intensity factor [30, 33]. A third interpretation [14, 24] concentrates on the way load is partitioned between the coating and the substrate for if the coating cracks, the load it was carrying must be redistributed. Because of constraint factors this
redistribution manifests itself as a stress concentration at the tip of the crack. At low temperatures the coating will be strong and hence any crack will have a large stress intensity at its tip. In contrast, when the coating is weak (high MPa
60O
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• CYCLES
Fig. 13. Maximum stress in the substrate v s . number of cycles to grow a coating crack to fail the single crystal test piece [14].
640 temperature) it carries little load anyway, hence the presence of a crack requires little load to be redistributed. Thus the stress intensity at the tip of a geometrically equivalent crack will be far smaller than when the coating was strong. The crack growth rate beyond the coating will therefore decrease with increasing temperature when the coating softens significantly. Although less explicitly stated than in ref. 14, a somewhat similar observation has been made [31] that non-propagating cracks behave as they do because they are "not sufficiently severe from a stress intensity standpoint to reduce the l07 uncoated strength of these superalloys". This is still a qualitative argument. Verification requires detailed mathematical modelling of the load repartitioning behaviour using constitutive equations to describe the coating's properties. Whilst this approach has been used to model TMF it does not yet seem to have been applied to the simpler case of isothermal fatigue. However, it can be seen that of the three factors identified above, crack initiation and propagation are linked through their common dependence on the mechanical properties of the coating. The third factor, the relative rate at which substrate and coating progress towards failure, is worth commenting on since it illustrates a difference between polycrystalline and single crystal materials. Several published results using straincontrolled LCF show that coatings have little effect on the lives of polycrystalline superalloys (see, for example, ref. 34). Polycrystalline superalloys have shorter lives than [001] oriented single crystals when fatigued under strain control because of the larger stresses induced by their greater Young's modulus. However, since the strain-stress range experienced by the coating is not dependent, to a first approximation, on the modulus of the substrate, its intrinsic life will be independent of substrate type. Hence it can be anticipated that coatings will affect the lives of polycrystalline superalloys less than those of single crystals or unidirectionally solidified superalloys, as is observed. This illustration emphasises the need to use care when interpreting test results on such "composite" systems.
4. Thermo-mechanical fatigue Simple experimental data on the TMF properties of coated and uncoated superalloys is fairly readily available, e.g. based on tests in hot gas
streams or fluidised beds. However, because of the somewhat empirical nature of the data, detailed analyses have not been forthcoming. Various interpretations have been advanced which offer semiquantitative rationales of the observed behaviour, e.g. the use of the DBTT curve. It is only recently that detailed attempts have been made to analyse and model the deformation and failure of the coating under TMF conditions. Various forms of TMF cycles can be postulated. Those most commonly employed are shown in Fig. 14 together with the effect of thermal expansion mismatch on the strain range for the coating under out-of-phase cycling. However, it is generally agreed that the most damaging cycle is that which imposes a tensile strain on the coating at low temperature (out-of-phase cycle): e.g. the cyclic life for NiCoCrAIY and aluminised PWA 1480 single crystals increased from ca. 500-1500 cycles (respectively) for out-of-phase TMF to ca. 10 500 cycles for in-phase TMF. It is this out-of-phase cycle which is the most relevant test vehicle in attempts to analyse the performance of coatings for gas turbine blading. One approach to the problem has been phenomenological. This treats the coating as being subject to the same descriptive rules that have been found useful in the analysis of structural alloys, e.g. inelastic strain range, strain range partitioning or hysteretic energy. However, because a coated component is a composite system where failure occurs as the consequence of a series of sequential events, the life of a coated component, Nf, has to be expressed as [11] Nf=Nc+N~c+Nsp
where N¢ is the cycles to initiate a crack through the coating, N~ is the cycles for coating crack to IN PHASE
z +
J
OUT OF PHASE
J
< cc
OUT OF PHASE +COATING MISMATCH ctc
BITHERMAL
+ IT -
>cxs\"
222"'-'" %---',_~co.,~G SUa~TR~
Fig. 14. Various forms of TMF cycling: in-phase, out-ofphase, bithermal and out-of-phasewith allowancefor coating expansion mismatch strain.
641
penetrate a small distance into the substrate and N~p is the cycles to propagate the substrate crack to failure. The best descriptive model for Arc was found to be that of tensile hysteretic energy [26]
O 1~
where f~ (back stress) and K (drag stress) are strain history dependent internal state variables describing kinematic and isotropic cyclic hardening respectively. Using this approach [26] for Arc it is possible to correlate isothermal overlay coating lives and coating lives under TMF conditions to within a factor of about 2.5 (Fig. 16). Determination of the N~ and N~p terms in the cycle life equation remains to be addressed [11, 26]. Likewise, the aluminide coating model still needs to be developed [26]. One of the objectives of any such work is to predict lives, but it also ought to be capable of indicating why different coatings behave differently. It was noted in the introduction that a
0600 L00
0
,oo yI DATA
-
C/3
~200 I o,"
/
/
427-1149"C
D [3
Y
/o)/ 102
.
102
,
,
103
10 ~"
ACTUAL LIFE
(CYCLES)
Fig. 16. Overlay coating life model correlation of isothermal life data and prediction of TMF lives [26].
major failing of an approach such as the DBTT route was its inability to explain why the relative TMF resistance of different coatings should be cycle- or blade-dependent [2]. A simplistic interpretation of the model would be that the "more ductile" coating would always be best. This is not found in practice, nor in TMF tests, although the data [37, 38] is as yet somewhat limited (Fig. 17). Here there appears to be a switch over in relative performance with the NiCoCrAIY PWA 286 being best at high strain ranges and the PWA 273 aluminide being (potentially) superior at low strain ranges. As yet no mechanistic reason for this difference appears to have been put forward, partly because of the lack of constitutive models for an aluminide coating. A somewhat different approach to the problem has placed more emphasis on the role of defects within the coating [4, 39] based on work with eb-pvd NiCoCrA1Y coatings on a variety of substrates. Once the extra strain induced in the coating by thermal expansion mismatch had been taken into consideration, coating cracking lives
~
PREDICTION ~OC) / O -
v
,~10 z
g
}--
z
0 -200
D TMF- 427 1038~C
1o
where A Wt is the tensile hysteretic energy, v is a complex time and temperature-dependent frequency term [35] and b and m are constants. The values of A Wt are derived from calculated hysteresis loops for the coating. These loops are, in turn, based on constitutive models for the coating. For PWA 286, NiCoCrA1Y, the best agreement between an experimental hysteresis loop and a constitutive model (Fig. 15) was found for the isotropic unified viscoplastic formulation [11, 36]
LCF
z.27760,927,038°C
N~ = C(A lYt) -bvm
O_
ISOTHERMAL
0k ~
J
OUCTP~F-P H A S E o
o;2 STRAIN
66
(%)
0.1
102
, L '
10 3
0
NICoCrAiY PWA 286
•
A l u m l n l d e PWA 273 , , __
,
10 ~
50"/. LOAD DROP LIFE (CYCLES)
Fig. 15. Comparison between experimental and calculated hysteresis loop (unified viscoplastic) for PWA 286 under TMF cycling [26].
Fig. 17. Total strain range v s . life for coated PWA 1480 single crystal for an out-of-phase TMF cycle: 427-1038 °C [37].
642 06
5. Concluding remarks
~z~ :E Lu I< P(z Dm
0/.
QZ)~O
SUBSTRATE 0 MAR-M200÷Hf
" 0 18 %
B 1900+Hf
- 0 18"/o
[] NX-188
~
o~ 10 2
J
,
t
~E~
~ L I 11103
- 0 29% ,
,
,
5x103
CYCLES TO CRACK INITIATION
Fig. 18. NiCoCrAIY coating cracking life at defects for an out-of-phase TMF cycle (427-1038°C) vs. mechanical + thermal mismatch strain [39].
could be correlated with the total strain range for out-of-phase cycling (Fig. 18). An important point arising from this, apart from the role of the expansion mismatch strain, is that initiation occurred at flaws. Constitutive models are based on "bulk" property data. A similar difficulty was noted in the isothermal fatigue section between "bulk" and in situ properties. This again raises the question of the importance of the surface condition and nature of structural flaws or weakness in such a high surface-to-volume "test piece" as an in situ coating. Whilst plasma-sprayed coatings will not have the same structural features as eb-pvd coatings, they are not metallurgically perfect products. It has been observed in the work on plasma-sprayed NiCoCrA1Y PWA 286 that crack initiation occurs at many sites over the coating, with cracks from many sites contributing to produce the failure. In contrast, failure in PWA 273-aluminised single-crystal PWA 1480 was dominated by long circumferential cracks [37]. The question of coating crack growth into the substrate has also been analysed [40], the coating crack being regarded as a geometrical extension of any (short) substrate crack, its effect being somewhat modified by differences in elastic constants and differential expansion. However, as was the case for the equivalent isothermal case, the coating crack cannot be treated purely as a geometric feature: its load-carrying capacity must be taken into account for this will affect AK at the crack tip. This will depend upon the temperature dependence of the coating's properties as well as the rate at which the TMF cycle is carried out and will greatly complicate any such analysis, especially for "slow" thermal transients. In addition to all these factors, environmental effects must be included, for TMF lives in vacuum are greater than those in air [29].
This brief survey of the mechanical properties of coatings and coated components has been very selective. Many factors have not been discussed, e.g. the effects of oxidation and corrosion, interdiffusion and ageing, or the effects of different coating techniques on the properties of MCrA1Y coatings. These are all additional effects which, whilst of importance, further complicate an already complex subject. Although it is clear that advances in understanding and component lifing are being made, there are many topics which are unresolved or simply unclear because of the lack of data, e.g. the relative merits in a high surfaceto-volume material such as a coating of a phenomenological approach based on "bulk" properties vs. an analysis based on crack initiation at surface flaws. However, any fracture mechanics type of approach is likely to be complex both because of the size of the defect but also because of the huge range of coating properties encompassed during a single TMF cycle. The answers to other important questions, such as the causes of the differences between MCrA1Y and aluminide coatings, is hampered by the lack of suitable data, in this case primarily because of experimental difficulties. Whilst deformation modelling has been actively pursued, the basic metallurgical side of the problem seems to have been somewhat overshadowed, e.g. there has been minimal work on crack initiation and growth. Although it is important to be able to life a particular coating or component, there is also the need to know how to modify it metallurgically in order to develop a better coating or improve deposition techniques. The increased effort and interest shown in this topic over recent years has brought into clearer focus some of the physical and metallurgical factors governing coating life. However, there is a great deal of work needing to be done before a sound quantitative understanding has been achieved.
Acknowledgment The author is grateful to the Directors of ERA Technology Ltd. for permission to publish this article.
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