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The oxidation and rumpling behavior of overlay B2 bond coats containing Pt, Pd, Cr and Hf
Surface & Coatings Technology 221 (2013) 13–21
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The oxidation and rumpling behavior of overlay B2 bond coats containing Pt, Pd, Cr and Hf R.W. Jackson a,⁎, D.M. Lipkin b, T.M. Pollock a a b
Materials Department, University of California, Santa Barbara, CA, USA GE Global Research, Niskayuana, NY, USA
a r t i c l e
i n f o
Article history: Received 18 August 2012 Accepted in revised form 5 January 2013 Available online 24 January 2013 Keywords: Oxidation Rumpling Spalling Palladium Platinum
a b s t r a c t The behavior of Hf-modified B2 nickel aluminide bond coatings, fabricated by ion-plasma deposition, has been investigated under thermal cycling conditions. The mass change, surface roughness, and microstructural evolution have been characterized. Variation in the Al, Cr, Pd and Pt concentrations did not affect the initial oxidation or rumpling behavior of these bond coats, but did influence oxide adherence as well as the surface roughness at failure. Both the Pt and, interestingly, the Pd-modified coatings displayed superior cyclic oxidation lifetimes compared to the state-of-the-art platinum aluminide coatings. The rate of surface roughening (rumpling) is discussed in terms of the B2 to L12 phase transformation and oxide spallation. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Nickel aluminide B2 bond coats are frequently applied to superalloy components used in gas turbine engines [1–3]. Owing to their high aluminum concentration, these coatings readily form alumina scales and are capable of maintaining excellent oxidation resistance in high-temperature environments for extended periods of time, even at coating thicknesses of less than 50 μm [4–7], as needed for many aero-engine applications. Oxide spallation, driven by stresses that develop due to the thermal expansion mismatch between the oxide and the superalloy, is a predominant degradation mechanism of these coated components [8,9]. Spallation occurs when the driving force for delamination, the stored elastic strain energy, exceeds the fracture toughness of the TGO/bond coat interface [10]. Because the CTE mismatch in the oxide–superalloy system is difficult to reduce, the resistance to spallation can therefore be enhanced by decreasing the growth rate of the TGO and by increasing the interfacial fracture toughness of the TGO/ bond coat interface. Bond coat composition strongly affects interfacial fracture toughness. The most pronounced effects result from impurities such as S, C, and P, which can have a markedly deleterious effect on adhesion even in parts-per-million (ppm) concentrations [11,12]. Sulfur in particular has been found to promote spallation by segregating to the TGO/bond coat interface where it decreases the work of adhesion and facilitates the formation of interfacial voids [13,14].
⁎ Corresponding author. Fax: +1 805 893 8486. E-mail address: [email protected] (R.W. Jackson). 0257-8972/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.surfcoat.2013.01.021
Reactive element additions, such as Y, Zr, and Hf, can greatly affect oxidation resistance in concentrations less than 1 at.%. A number of mechanisms have been postulated to explain the improved TGO life, including gettering of sulfur [15,16], formation of oxide protrusions into the bond coat that impede crack propagation [17], slowing oxide growth by reduction of Al diffusion in the Al2O3 TGO [18], and by segregating to the TGO/bond coat interface limiting sulfur segregation [19]. The relative importance of each of these mechanisms is system-dependent and is still a subject of active research and discussion. B2-base bond coats are also frequently alloyed with platinum group metals (PGM), in concentrations ranging from 5 to 20 at.% [20–23]. The most common PGM addition is Pt. The state-of-the-art Pt-modified aluminide coatings contain approximately 7 at.% Pt [24]. Pt has been suggested to improve the oxidation resistance of aluminide coatings via a number of mechanisms. First, Pt lowers the activity of Al in B2 NiAl. As Pt accumulates near the TGO, due to the selective oxidation of other elements, Al activity gradient is increased, increasing the flux of Al to the TGO/bond coat interface promoting alumina scale formation [25]. Pt additions also improve scale adhesion by inhibiting S segregation to the scale/alloy interface [12,26]. As a result, the thermal cycling lifetimes of Pt-modified aluminide coatings are typically longer than simple nickel aluminide coatings under thermal cycling conditions. As an alternative to Pt, other PGM alloying additions have been investigated [23,27,28]. Of particular interest is Pd, which, while sharing many chemical similarities to Pt, is significantly less expensive on an atom-for-atom basis. However, the degree to which Pd additions improve oxidation resistance is a subject of debate [21,28–31]. The lifetime of Pd-modified coatings are often found to be greater
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R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
Table 1 Nominal bond coat compositions. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
Nominal composition at.% Ni
Al
Cr
Hf
60 56 55 60
34 33 34 23
6 6 6 12
0.3 0.3 0.3 0.3
Pd
Pt
5 5 5
[42]. However, the mechanism by which Hf inhibits rumpling has not been clearly demonstrated. In this work, Hf-modified B2 bond coats with varying Pt, Pd, and Cr additions were thermally cycled to 1150 °C and the oxidation and rumpling behaviors were investigated. The influence of the transformation from B2 to L12 on the system evolution is studied in detail. The performance of these coatings is compared with the stateof-the-art platinum aluminide coatings. 2. Materials and experimental procedure
than unalloyed nickel aluminides, but less than Pt-modified coatings. Also, it is not known how Pd affects the Al activity in B2 NiAl, or S segregation. Even the phase equilibrium in the Ni–Al–Pd system is poorly understood. Rumpling is another phenomenon that can degrade bond coats [32–35]. The low creep resistance of B2 bond coats makes them particularly susceptible to morphological evolution when stressed [32,36–38]. These stresses can develop at operating temperature, as a result of TGO growth stresses, or interdiffusion-induced stresses in the intermetallic. During thermal cycling, stresses also develop due to differences in the coefficient of thermal expansion (CTE) between the TGO, the bond coat, and the superalloy, as well as phase changes occurring during heating and cooling. The relative importance of each of these factors depends upon the details of the thermal history and coating composition. Most B2 coatings have been found to rumple when cycled to maximum temperatures above 1100 °C, with Pt and Al concentrations not strongly affecting the rumpling rate [39–41]. Hf additions, on the other hand, have been observed to decrease the rate of the rumpling
70
a
Ni
Concentration (atomic %)
60 50 40 30 20
Al Cr
10
Pd
0 0
10
20
30
40
50
60
70
Depth (µm) 10
b
Bond Coat
Superalloy 1.5
8 7
Mass Change (mg/cm2)
Sulfur Concentration (appm)
9
Single crystal René N5 (Ni–7.6Co–8Cr–0.9Mo–1.6W–1Re–2.2Ta– 13.7Al–0.05Hf–0.75C–0.02B–0.01Y, at.%) was used as the substrate. Cylindrical disk coupons, 12 mm in diameter and 2 mm in thickness, were coated with four bond coat compositions using ion plasma deposition (IPD) over the entire surface to an average thickness of approximately 50 μm [43]. The coating compositions, measured by X-ray fluorescence (XRF), are presented in Table 1. Three of the coatings, designated 5Pt, 5Pd, and + 5Ni, contained the same concentration of Al, Cr, and Hf with additions of 5 at.% of the metals, Pt, Pd, and Ni, respectively. The fourth coating contained 5 at.% Pd, with less Al but more Cr. Using thermodynamic databases, the Ni and Al contents were adjusted to maintain the B2-NiAl structure. Following deposition, the coatings were heat-treated for 4 h at 1079 °C under vacuum, and then lightly grit blasted (30 psi, 220 grit Al2O3), imparting a surface roughness of approximately 2 μm. The composition depth profile was measured using glow discharge mass spectroscopy, (GDMS), by Evans Analytical Group (Fig. 1). The S concentration profile, also measured by GDMS, is less than 1 atomic ppm (appm) in the bond coat, and less than 5 appm in the superalloy (Fig. 1). The thermal cycle used to oxidize the coatings consisted of a 45 minute dwell period at 1150 °C, with heating and cooling rates of approximately 200 °C/min. Two samples of each bond coat composition, + 5Ni, 5Pd, 5Pt, and 5Pd–12Cr–23Al, were cycled until a net mass loss had occurred, which was defined as failure. A third set of samples underwent 500 cycles. Surface and cross-sectional microstructural analysis was performed using the scanning electron microscope (SEM). The surface morphology of the bond coat specimens was analyzed with an optical surface profilometer, using vertical scanning interferometry, as described elsewhere [44]. Surface height measurements, with a vertical resolution better than 0.1 μm, were scanned over an area of 0.46 × 0.62 mm2 at 10× magnification with a lateral resolution of 1.0 μm, and over an area 0.92× 1.24 mm2, at 5× magnification with a lateral resolution of 1.9 μm. Five measurements were recorded at each magnification. The surface topology was characterized by the root-mean-squared (RMS) roughness, Rq (Eq. (1)), where zi is a measured surface height, z ̅ is the
6 5 4 3 2 1
+5Ni
1.0
5Pd
5Pt
0.5 Pt-Aluminide*[45] 5Pd-12Cr-23Al
0 0
10
20
30
40
50
60
70
Depth (µm)
0.0
0
500
1000
1500
2000
Cycles Fig. 1. The bulk composition profile (a) and sulfur concentration (b) of the as heat-treated 5Pd-12Cr-23Al, measured by GDMS.
Fig. 2. Mass change per unit area of the bond coats thermally cycled at 1150 °C.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
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Table 2 Number of cycles to maximum mass change and failure, as defined by zero crossing. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
# of cycles Max. mass gain
Failure
267 400 1050 267
975 1050 1775 850
average height, and n is the number of locations where the height was measured in each scan. vffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi u n u1 X 2 Rq ¼ t ðz −z Þ : n i¼1 i
ð1Þ
The RMS roughness can be related to the scale of the surface undulation, by envisioning a sinusoidal surface of the form 2πx 2πy sin zðx; yÞ ¼ A sin L L
ð2Þ
where A is the amplitude, L is the wavelength, and A = 2Rq. 3. Results Bond coat specimens were thermally cycled to a maximum temperature of 1150 °C, and the specific mass change was measured
Fig. 4. Secondary electron SEM image of 5Pt coating after 500 cycles showing cracks in the TGO at the apex of coating ridges.
periodically (Fig. 2). The mass change of grit-blasted, Pt-aluminide coated René N5, reported by Tolpygo et al. [45], is overlaid for comparison. During cyclic oxidation there was an initial period of mass gain, during which the rate of mass gain decreased with cycle number until a maximum was reached, followed by mass loss, during which the rate of mass loss increased with cycle number. The rate of mass change during the initial 100 cycles was similar for all samples tested. Table 2 lists the average number of cycles to maximum weight gain at which the rate of specific mass change became negative, as well as the average number of cycles to failure, as defined by crossing zero.
Fig. 3. Plan view back-scatter SEM micrograph of the coatings after 400 cycles: showing regions of TGO spallation and grain boundary ridges (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt.
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The oxide scale formed on all of the alloys was primarily composed of Al2O3 with fine entrained HfO2 particles, typically 0.1–1 μm in diameter. Surface backscatter SEM images of each bond coat after 400 cycles are presented in Fig. 3, where the dark-contrast phase is Al2O3 and the small bright phases are HfO2. The bright HfO2 particles can be observed on the surface of the 5Pt alloy after 500 cycles (Fig. 4). Cracks in the scale were observed after 500 cycles, with the majority of these cracks residing at the apex of surface undulations (Fig. 4). Oxide spallation was observed emanating from these cracks after 200 cycles in the 5Pd–12Cr–23Al and +5Ni, after 250 cycles in the 5Pd, and after 600 cycles in the 5Pt. Areas of exposed bond coat are evident as the large bright regions in Fig. 5. Areas lacking HfO2 particles, that are roughly the same size as the exposed metal regions, are interpreted to be re-oxidized spalled regions. Higher magnification micrographs of the bond coat surface below spalled oxide (Fig. 5) show imprints of the inwardly growing Al2O3 grains. Oxide pegs, areas of the TGO that grow into the alloy along Hf-rich stringers are evident in the exposed bond coat underneath TGO spalls. Cross-sectional backscattered SEM images after 500 cycles (Fig. 6) confirm that the alloys form an Al2O3 scale with a high density of HfO2 particles near the TGO surface, with occasional larger HfO2 particles throughout the scale. Oxide scales where HfO2 is present at the surface indicated regions that have not spalled and re-oxidized, and the scale thickness in these regions is similar for all alloys, 4–6 μm. The constant scale thickness in non-spalled regions is in agreement with the scale mass change during the initial stages of oxidation.
The frequency of oxide spallation increases concurrently with the decrease in specific mass change. The amount of spalled area is largest for the 5Pd–12Cr–23Al bond coat and smallest for the 5Pt bond coat. Spalled regions are apparent in cross-section (Fig. 7) as regions of bare metal as well Al2O3 absent HfO2, the latter indicative of re-grown oxide. Cross-sectional micrographs of coatings cycled to failure (Fig. 7) show that the oxide scale is still primarily composed of Al2O3 although more of the scale has spalled. Additionally, only a small volume fraction of nickel oxide is observed in the scale. All of the bond coats extensively rumpled during thermal cycling. This is illustrated by the secondary electron imaged surfaces in Figs. 8 and 9. The extent of rumpling can also be observed in the crosssections in Figs. 6 and 7 as well as in the surface contour maps in Fig. 9 that show the 5Pt coating after grit-blasting and after 250 cycles. The RMS roughness as a function of cycle number is presented in Fig. 10. The rate of rumpling was similar for all of the bond coats for the first 100 cycles. The 5Pd–12Cr–23Al coating rumpled the least overall, and little change in the surface roughness was observed after 200 cycles. Rumpling ceased in the 5Pd bond coat after 350 cycles, and the + 5Ni sample experienced no further rumpling after 500 cycles. The 5Pt samples continued rumpling through 1500 cycles, developing with the highest rumpling amplitude of all of the bond coats. The evolution of surface roughening can best be discerned by dividing the lifetime of each sample into three regimes. At the onset of thermal cycling, the roughness of all compositions increases significantly and at a nearly identical rate. Subsequently, the bond coats rumple at different relative rates until a plateau roughness is reached prior to failure. The phase evolution within the bond coats was analyzed by X-ray diffraction (XRD). The initial microstructure of all the coatings consisted solely of B2 NiAl, with the exception of 5Pd–12Cr–23Al, which contained a small fraction of L12 Ni3Al. During cooling in the first thermal cycle, the B2 phase in all samples underwent a martensitic transformation to tetragonal L10 and the monoclinic 7R [46]. During high-temperature exposure, Al is lost from the bond coat due to TGO growth and interdiffusion with the substrate. As a result of Al depletion, the B2 (7R and L10) NiAl phase transforms to the L12 Ni3Al phase. Table 3 lists the time at which L12 is first detected, as well as the point at which the diffraction of L10 is no longer detected. Fig. 11 shows the XRD scans for 5Pd, illustrating the growth of the L12 (111) reflection and the corresponding decrease of the 7R (113) and L10 (200) reflections. 4. Discussion While the initial oxidation and rumpling behaviors of the four alloys were similar, a large variation in the cyclic oxidation lifetime and surface roughness at failure was observed. The discussion below focuses on identifying which compositional factors played a role in the stratification of these phenomena through thermal cycling. 4.1. Cyclic oxidation behavior
Fig. 5. Back-scatter SEM images of +5Ni after 500 cycles. (Top) Spalled regions in which the metallic bond coat surface is visible are outlined in red and re-oxidized regions that do not contain HfO2 are outlined in blue. (Bottom) Higher-magnification images of a spalled area and (inset) HfO2 particles in the TGO, outlined in blue.
The net mass change and oxide morphology during the first 100 cycles are similar for all alloys. Limited spinel formation and no oxide spallation were observed. These results demonstrate that the Al and Cr concentrations in all of the bond coats were sufficient to grow alumina at 1150 °C and that PGM additions are not necessary to prevent the growth of transient oxides, as is the case in bond coats with lower Al and Cr concentrations [6,20,25]. Many alloys exposed to high-temperature thermal cycling fail when their ability to form a protective oxide through selective oxidation is compromised due to compositional change via interdiffusion and/or loss of Al to the oxide [8]. However, alumina growth continued
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
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Fig. 6. Cross-sectional back-scatter SEM images of the (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt bond coats after 500 cycles, showing the rumpling under the TGO scales. Higher-magnification (insets) shows the presence of embedded Hf-rich particles in the TGO.
throughout the thermal cycling test for all four alloys, even as localized spallation and, ultimately, net mass loss occurred. Overlay coatings are not in thermodynamic equilibrium with the substrate, so a flux of elements into the superalloy during thermal cycling is expected. Pt, which decreases the activity of Al in Ni–Al–Pt alloys [47], is expected to decrease the flux of Al into the superalloy by decreasing the Al chemical potential gradient. Pd could be expected to have a similar affect on the Al activity and its interdiffusion behavior. In a parallel set of coatings fabricated by the same approach as described here, the concentration of the Pd in the coating was found to fall to approximately 25% of its original level, while the Pt only fell to about 45% of its original level after 1800 cycles at 1150 °C [Adharapurapu, unpublished research]. Therefore, a lower Al concentration at the TGO interface might be expected in the Pd-containing coatings compared to the Pt-containing coatings [23]. However, because Al2O3 is formed throughout cycling it does
not appear that Al influences the ability of the coatings to reform the protective oxide, within the range of composition employed here. As shown in Fig. 4, the major difference in the response of the four coatings is in the frequency of the spalling. The Pt-containing coating experiences a lower incidence of spalls throughout cycling. Previous high-temperature studies of the mechanical properties of B2 coatings [23] suggest that the differences are dominated by interfacial effects. However, isolating the role of alloying additions on the oxidation of B2 bond coats during thermal cycling is confounded by changes in properties as well as interdiffusion with the substrate. The effect of composition on the Al2O3/bondcoat interface has been the subject of many experimental and theoretical investigations. When the toughness has been directly measured and where the work of adhesion has been calculated through sessile drop tests, it has been found that bulk alloying additions of Cr do not markedly improve
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R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
Fig. 7. Cross-sectional back-scatter SEM images of the bond coats cycled to failure (A) +5Ni, 850 cycles; (B) 5Pd, 1000 cycles; (C) 5Pd-12Cr-23Al, 850 cycles and (D) 5Pt, 1750 cycles.
adhesion [48–50], while Pt additions can increase the work of adhesion by approximately 20% [51]. Smaller improvements can be observed for Y [49], and Al. The effect of composition on interfacial adhesion has been studied more systematically through the use of density-functional theory [52]. These results are in general agreement with the experimental work, and suggest that Cr, Al, Pt, and Hf additions do not markedly increase the work of adhesion between nickel and alumina [19,52]. The metal/Al2O3 work of separation typically falls in the range of 2.5–3.5 J/m 2 and depends on interfacial stoichiometry and the orientational relationship as much as the composition. The influence that alloying Ni with Pd has on the Al2O3/Ni adhesion has not be reported on in the literature, but the work of separation between Pd and Al2O3, is less than Pt, Ni, and Al respectively [52]. Unlike Al, Cr, Pt, Pd, and Hf, sulfur is known to have a strong effect on interfacial toughness [12,15,19,52–55]. There are two primary modes by which sulfur segregation degrades interfacial toughness:
facilitation of the growth of interface voids [5,11], and, the increase in alloy–oxide interfacial energy [53]. Both of these mechanisms can be affected by alloying additions. However, the total S concentration in the bond-coated superalloys is low, less than 1 atomic ppm in the bond coat and 4 ppm in the superalloy (Fig. 1). Nevertheless, with increased exposure time at 1150 °C, S is expected to diffuse toward the TGO due to the large heat of segregation [13,14]. It is therefore important to consider how the alloying additions affect the S segregation. Reactive elements, such as Hf, are known to mitigate the detrimental effect that S has on adherence. HfS2 has a large negative free energy of formation [56] and therefore Hf dissolved in the bond coat will limit the S activity in the alloy. Additionally, Hf segregates to the alloy/alumina interface and because the heat of segregation is larger than that of sulfur, the concentration of sulfur at the interface is limited [19]. However, all of the bond coats investigated contain
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
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Fig. 8. Plan view secondary electron SEM micrographs of the coatings (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt after 400 cycles.
Hf, and a large range in spallation resistance was observed. It is of interest to discuss how Cr, Al, Pt and Pd additions may affect S segregation, and, in turn, adherence. Both theoretical and experimental investigations have shown that Cr promotes the interfacial segregation of S in Ni–Al2O3 systems through a Cr–S co-segregation process. Accordingly, the interface strength of Ni–Al2O3 systems in which sulfur is present is
markedly reduced by Cr additions. Further, Hou has shown that Cr additions can overwhelm the beneficial effects of reactive elements [14,26]. Therefore, the co-segregation of Cr and S to the BC/TGO interface is a plausible explanation for the poor spallation resistance of the 5Pd–12Cr–23Al coating, with respect to the other coatings investigated. Pt additions, widely used in commercial bond coats, have been shown to mitigate the effects of S by reducing the formation of
Fig. 9. Plan view secondary electron SEM micrograph of the 5Pt coating (A) after 0 and (B) 250 thermal cycles and contour map of 5Pt after (C) 0 and (C) 250 cycles showing the topography of the rumpled surface.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
Roughness Rq (μm)
15
5Pt
10 +5Ni 5Pd
5
0
5Pd-12Cr-23Al
0
200
400
600
800
1000
1200
Cycles Fig. 10. RMS roughness of the bond coats thermally cycled at 1150 °C.
interfacial voids and limiting the amount of interfacial S segregation. Accordingly, the 5Pt bond coat showed the greatest spallation resistance. However, the relative importance of Pt and Hf is illustrated when the 5Pt bond coat is compared to the CVD platinum aluminide (Fig. 2) which had a thermal cycling lifetime comparable to the +5Ni sample. Replacing Pt with Pd resulted in a marked decrease in the spallation resistance, but the 5Pd bond coat was still superior to the +5Ni bond coat. The paucity of reported investigations of Pd-modified nickel aluminides system makes it difficult to infer the mechanistic root of the benefit imparted by Pd. Oquad and Monceasu have reported that interfacial voids form when B2 (Ni,Pd)–Al are oxidized at 1100 °C, similar to binary B2 NiAl and in contrast to B2 (Ni,Pt)–Al. However, few voids were observed in the 5Pd bond coat, as is commonly the case for Hf-modified bond coats. It is reasonable to speculate that Pd, like Pt segregates to the metal–Al2O3 interface, inhibiting S segregation and further work to examine this hypothesis would be useful. Finally, it is noted that the concentration of Hf in the coatings was relatively high, and can be considered to be ‘over-doped’ due to the amount of HfO2 present in the TGO. A decrease in the HfO2 concentration from 0.3 to 0.05–0.1 at.% has been shown to increase the cyclic oxidation behavior of NiAl alloys [5] and could improve the oxidation resistance of the coatings investigated in this study.
4.2. Rumpling All of the bond coats in this investigation were found to rumple extensively. The rate at which the surface roughened was similar for all of the alloys during the first 100 cycles. This rumpling behavior is typical of B2 bond coats on René N5 superalloys [36,39]. The primary factors that drive rumpling have been identified as growth strains in the TGO and thermal expansion mismatch strains in the superalloy–bond coat–TGO system [57,58]. These strains should not initially vary significantly among the compositions investigated, given the similarity in the initial oxidation rate and the minor effect that the composition has on the CTE of B2 NiAl bond coats [4], consistent
with the similarity of rumpling rates in all coatings early in the cyclic life. As thermal cycling continues, several phenomena that may affect rumpling are occurring: the B2 to L12 phase transformation, bond coat–superalloy interdiffusion, and oxide spallation. As the Al content in the bond coat decreases, the B2 structure transforms into the L12 phase. This process is expected to affect the rate of rumpling in several ways. First, the B2 phase transforms via a martensitic transformation to L10 during cooling [59]. The volumetric expansion associated with this phase transformation induces stresses that promote rumpling, as illustrated by Balint and Hutchinson [57]. The reduction of the volume fraction of B2 should reduce the amount of strain imposed by the martensitic transformation. The CTE mismatch strain between the bond coat and the superalloy will also change as the phase transformation proceeds. Haynes et al. [60] have reported the average CTE of several Ni–Al alloys including of a B2 alloy, Ni–39Al–5Pt, a L12 alloy, Ni–25Al–Hf, and the two-phase γ/γ′ (L12/FCC) superalloy René N5 (16.0, 15.6, 17.2 ppm K − 1 respectively at 1150°). Therefore, as the B2–L12 transformation proceeds, the CTE misfit strain between the bond coat and the superalloy is expected to increase. If further Al depletion occurs the L12 phase will transform to FCC γ-NiAlCr solid solution. The CTE of γ phase, 18–19 ppm K − 1 [60] is higher the L12, B2 phases and René N5, which will decrease the CTE mismatch strain. Finally, the L12–B2 phase transformation will change the creep behavior of the bond coat. While both B2 and L12 are soft at 1150 °C [61,62], the flow stress of Ni3Al is nearly twice that of NiAl at 800 °C, 180 MPa and 95 MPa respectively [63,64]. Additionally, while the flow stress of B2 decreases monotonically from 800 °C to room temperature, L12 Ni3Al exhibits anomalous behavior with a maximum flow stress around 600 °C [65]. The non-linear temperature dependence of L12 flow stress and the René N5 CTE complicate the analysis of the effect that the B2/L12 volume fraction has on the rumpling rate. The B2 to L12 phase transformation is driven by Al depletion, resulting from Al2O3 growth and interdiffusion with the superalloy, and the volume fraction of the L12 phase increases with time. The rate of L12 growth is greatest in the 5Pd–12Cr–23Al bond coat due to spallation and low initial Al concentration, while the transformation occurs most slowly in the 5Pt bond coat, which exhibits the least oxide spallation and Al back-diffusion [47]. However, the extent of rumpling is also greater for the 5Pt bond coat. The amplitude of the surface undulations may also be affected by oxide spallation, in addition to the thermo-mechanical stress induced deformation of the bond coat. The majority of spallation events result from cracks which form at the apex of surface undulations, as
Intensity (A.U.)
20
Table 3 Number of cycles at which the L12 phase first was detected and after which L10 was no longer detected by X-ray diffraction. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
# of cycles L12 start
No NiAl
10 10 50 0
400 600 600 400
40
42
44
46
48
2θ
50
52
54
(o)
Fig. 11. X-ray diffraction pattern of 5Pd after 50 and 500 cycles showing the increase in intensity of L12 diffractions peaks corresponding to the increasing volume fraction.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
illustrated in Fig. 5. These observations are in agreement with the analysis of Hutchinson [66] who calculated the energy release rate for compressed films on curved surfaces, showing that the driving force for delamination is greatest on a convex surface. The selective removal of material at the top of undulations will have the tendency to reduce the overall surface roughness. Accordingly, the order in which spallation was observed in the bond coats, first 5Pd–12Cr– 23Al, followed by + 5Ni, 5Pd, and 5Pt, is the same as the order of observed rumpling rate reduction. The plateau in roughness may therefore be due to a combined effect of spallation and phase transformation. 5. Conclusions A series of Hf-modified B2 bond coats were thermally cycled to a maximum temperature of 1150 °C and the oxidation and rumpling behaviors were characterized. All of the bond coats formed Al2O3 scales. The Pt + Hf containing bond coat was the most spallation resistant and had the longest cyclic oxidation lifetime. The 5Pd–12Cr–23Al coating, which had the lowest Al and highest Cr concentration was the most prone to spallation. The 5Pd and + 5Ni bond coats had similar cyclic oxidation lifetimes as the state-of-the-art, grit-blasted Pt-modified diffusion aluminide bond coats, while the 5Pt lasted nearly twice as many cycles. All of the bond coats investigated were found to rumple. The rate of rumpling decreased with continued thermal cycling. The relative reduction in rumpling rates may be attributed to three factors: (1) the increased creep resistance of the L12 phase, which evolved during high-temperature exposure; (2) the removal of B2 phase which undergoes martensitic phase transformation upon cooling and (3) oxide spallation that, occurring preferentially at convex regions of the rumpled structure, dampened the rumpling amplitude. Acknowledgments The authors are grateful for the support of a NSF-Goali program, DMR Grant DMR-0605700. The authors acknowledge R. DiDomizio with coating processing, V. Dheeradhada for the XRF bond coat composition measurements, and R. Adharapurapu and C. Levi for the helpful discussions. References [1] G.W. Goward, D.H. Boone, Oxid. Met. 3 (1971) 475. [2] A.G. Evans, D.R. Mumm, J.W. Hutchinson, G.H. Meier, F.S. Pettit, Prog. Mater. Sci. 46 (2001) 505. [3] I.T. Spitsberg, D.R. Mumm, A.G. Evans, Mater. Sci. Eng., A 384 (2005) 176. [4] J.A. Haynes, B.A. Pint, Y. Zhang, I.G. Wright, Surf. Coat. Technol. 204 (2009) 816. [5] B.A. Pint, K.L. More, I.G. Wright, Oxid. Met. 59 (2003) 257. [6] C.S. Giggins, F.S. Pettit, J. Electrochem. Soc. 118 (1971) 1782. [7] F.S. Pettit, Trans. Metall. Soc. AIME 239 (1967) 1296. [8] G.H. Meier, Mater. Sci. Eng., A 120 (1989) 1. [9] H.E. Evans, R.C. Lobb, Curr. Sci. 24 (1984) 209. [10] A.G. Evans, G.B. Crumley, R.E. Demaray, Oxid. Met. 20 (1983) 193.
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[11] P.Y. Hou, K. Priimak, Oxid. Met. 63 (2005) 113. [12] J.A. Haynes, B.A. Pint, K.L. More, Y. Zhang, I.G. Wright, Oxid. Met. 58 (2002) 513. [13] W. Zhang, J.R. Smith, X. Wang, A.G. Evans, Phys. Rev. B 67 (2003), (245414-245411-245412). [14] P.Y. Hou, Annu. Rev. Mater. Res. 38 (2008) 275. [15] A.W. Funkenbusch, J.G. Smeggil, N.S. Bornstein, Metall. Trans. A 16 (1985) 1164. [16] J.G. Smeggil, Mater. Sci. Eng., A 87 (1987) 261. [17] J. Stringer, Mater. Sci. Eng., A 120–1 (1989) 129. [18] B.A. Pint, Oxid. Met. 45 (1996) 1. [19] Y. Jiang, J.R. Smith, J. Mater. Sci. 44 (2009) 1734. [20] E.J. Felten, Oxid. Met. 10 (1976) 23. [21] S. Alperine, P. Steinmetz, A. Friant-Constantini, P. Josso, Surf. Coat. Technol. 43 (1990). [22] B. Tryon, F. Cao, K.S. Murphy, C.G. Levi, T.M. Pollock, JOM 58 (2006) 53. [23] T.M. Pollock, D. Widrevitz, R. Pong, F. Cao, B. Tryon, Mater. Res. Soc. 1128 (2008) 343. [24] B.M. Warnes, D.C. Punola, Surf. Coat. Technol. 94–95 (1997) 1. [25] B. Gleeson, N. Mu, S. Hayashi, J. Mater. Sc. 44 (2009) 1704. [26] P.Y. Hou, T. Izumi, B. Gleeson, Oxid. Met. 72 (2009) 109. [27] B. Tryon, K.S. Murphy, J.Y. Yang, C.G. Levi, T.M. Pollock, Surf. Coat. Technol. 202 (2007) 349. [28] D. Oquab, D. Monceau, Scr. Mater. 44 (2001) 2741. [29] S.J. Hong, G.H. Hwang, W.K. Han, S.G. Kang, Intermetallics 17 (2009) 381. [30] R. Swadzba, M. Hetmanczyk, M. Zozanska, W. Bartosz, L. Swadzba, Surf. Coat. Technol. 206 (2011) 1538. [31] M.J. Li, X.F. Sun, H.R. Guan, X.X. Jiang, Z.Q. Hu, Surf. Coat. Technol. 167 (2003) 106. [32] P. Deb, D.H. Boone, T.F. Manley, J. Vac. Sci. Technol. 5 (1987) 3366. [33] V.K. Tolpygo, D.R. Clarke, Acta Mater. 48 (2000) 3283. [34] C. Mercer, D. Hovis, A.H. Heuer, T. Tominatsu, Y. Kagawa, A.G. Evans, Surf. Coat. Technol. 202 (2008) 4915. [35] C. Mercer, Y. Kagawa, T. Tominatsu, D. Hovis, T.M. Pollock, Surf. Coat. Technol. 205 (2011) 3066. [36] V.K. Tolpygo, D.R. Clarke, Acta Mater. 52 (2004) 5115. [37] V.K. Tolpygo, D.R. Clarke, Acta Mater. 52 (2004) 5129. [38] V.K. Tolpygo, D.R. Clarke, Surf. Coat. Technol. 203 (2009) 3278. [39] V.K. Tolpygo, D.R. Clarke, Acta Mater. 52 (2004) 615. [40] V.K. Tolpygo, Surf. Coat. Technol. 202 (2007) 617. [41] B. Hazel, J. Rigney, M. Gorman, B. Boutwell, R. Darolia, Proceedings of the International Symposium on Superalloys, 2008, p. 753. [42] V.K. Tolpygo, K.S. Murphy, D.R. Clarke, Acta Mater. 56 (2008) 489. [43] J.-C. Zhao, D.M. Lipkin, US Patent 6 (2005) 964,791. [44] V.K. Tolpygo, D.R. Clarke, Scr. Mater. 57 (2007) 563. [45] V.K. Tolpygo, D.R. Clarke, K.S. Murphy, Metall. Mater. Trans. A 32 (2001) 1467. [46] R. Kainuma, H. Ohtani, K. Ishida, Metall. Mater. Trans. A 27 (1996) 2445. [47] E. Copland, J. Phase Equilib. Diffus. 28 (2007) 38. [48] F.G. Gaudette, S. Suresh, A.G. Evans, G. Dehm, M. Ruhle, Acta Mater. 45 (1997) 3503. [49] R.M. Crispin, M. Nicholas, J. Mater. Sci. 11 (1976) 17. [50] C.A. Calow, P.D. Bayer, I.T. Porter, J. Mater. Sci. 6 (1971) 150. [51] A. Gauffier, E. Saiz, A.P. Tomsia, P.Y. Hou, J. Mater. Sci. 42 (2007) 9524. [52] H. Li, W. Zhang, J.R. Smith, Phys. Status Solidi A 208 (2011) 1166. [53] F.G. Gaudette, S. Suresh, A.G. Evans, Metall. Mater. Trans. A 31A (2000) 1977. [54] J.L. Smialek, JOM (2000) 22. [55] W. Zhang, J.R. Smith, X. Wang, A.G. Evans, Phys. Rev. B 67 (2003) 245414. [56] S.R. Shatynski, Oxid. Met. 11 (1977) 307. [57] D.S. Balint, J.W. Hutchinson, Acta Mater. 51 (2003) 3965. [58] D.S. Balint, T. Xu, J.W. Hutchinson, A.G. Evans, Acta Mater. 54 (2006) 1815. [59] R.J. Thompson, J.-C. Zhao, K.J. Hemker, Intermetallics 18 (2010) 796. [60] J.A. Haynes, B.A. Pint, W.D. Porter, I.G. Wright, Mater. High Temp. 21 (2004) 87. [61] D. Pan, M.W. Chen, P. Wright, K.K.J. Hemker, Acta Mater. 51 (2003) 2215. [62] K.J. Hemker, B.G. Mendis, C. Eberl, Mater. Sci. Eng., A 483 (2008) 727. [63] D.P. Pope, S.S. Ezz, Inter. Met. Rev. 29 (1984) 136. [64] K.R. Forbes, U. Glatzel, R. Darolia, W.D. Nix, Metall. Mater. Trans. A 27 (1996) 1229. [65] K.J. Hemker, M.J. Mills, W.D. Nix, Acta Metall. Mater. 39 (1991) 1901. [66] J.W. Hutchinson, J. Mech. Phys. Solid 49 (2001) 1847.
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Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
The oxidation and rumpling behavior of overlay B2 bond coats containing Pt, Pd, Cr and Hf R.W. Jackson a,⁎, D.M. Lipkin b, T.M. Pollock a a b
Materials Department, University of California, Santa Barbara, CA, USA GE Global Research, Niskayuana, NY, USA
a r t i c l e
i n f o
Article history: Received 18 August 2012 Accepted in revised form 5 January 2013 Available online 24 January 2013 Keywords: Oxidation Rumpling Spalling Palladium Platinum
a b s t r a c t The behavior of Hf-modified B2 nickel aluminide bond coatings, fabricated by ion-plasma deposition, has been investigated under thermal cycling conditions. The mass change, surface roughness, and microstructural evolution have been characterized. Variation in the Al, Cr, Pd and Pt concentrations did not affect the initial oxidation or rumpling behavior of these bond coats, but did influence oxide adherence as well as the surface roughness at failure. Both the Pt and, interestingly, the Pd-modified coatings displayed superior cyclic oxidation lifetimes compared to the state-of-the-art platinum aluminide coatings. The rate of surface roughening (rumpling) is discussed in terms of the B2 to L12 phase transformation and oxide spallation. © 2013 Elsevier B.V. All rights reserved.
1. Introduction Nickel aluminide B2 bond coats are frequently applied to superalloy components used in gas turbine engines [1–3]. Owing to their high aluminum concentration, these coatings readily form alumina scales and are capable of maintaining excellent oxidation resistance in high-temperature environments for extended periods of time, even at coating thicknesses of less than 50 μm [4–7], as needed for many aero-engine applications. Oxide spallation, driven by stresses that develop due to the thermal expansion mismatch between the oxide and the superalloy, is a predominant degradation mechanism of these coated components [8,9]. Spallation occurs when the driving force for delamination, the stored elastic strain energy, exceeds the fracture toughness of the TGO/bond coat interface [10]. Because the CTE mismatch in the oxide–superalloy system is difficult to reduce, the resistance to spallation can therefore be enhanced by decreasing the growth rate of the TGO and by increasing the interfacial fracture toughness of the TGO/ bond coat interface. Bond coat composition strongly affects interfacial fracture toughness. The most pronounced effects result from impurities such as S, C, and P, which can have a markedly deleterious effect on adhesion even in parts-per-million (ppm) concentrations [11,12]. Sulfur in particular has been found to promote spallation by segregating to the TGO/bond coat interface where it decreases the work of adhesion and facilitates the formation of interfacial voids [13,14].
⁎ Corresponding author. Fax: +1 805 893 8486. E-mail address: [email protected] (R.W. Jackson). 0257-8972/$ – see front matter © 2013 Elsevier B.V. All rights reserved. http://dx.doi.org/10.1016/j.surfcoat.2013.01.021
Reactive element additions, such as Y, Zr, and Hf, can greatly affect oxidation resistance in concentrations less than 1 at.%. A number of mechanisms have been postulated to explain the improved TGO life, including gettering of sulfur [15,16], formation of oxide protrusions into the bond coat that impede crack propagation [17], slowing oxide growth by reduction of Al diffusion in the Al2O3 TGO [18], and by segregating to the TGO/bond coat interface limiting sulfur segregation [19]. The relative importance of each of these mechanisms is system-dependent and is still a subject of active research and discussion. B2-base bond coats are also frequently alloyed with platinum group metals (PGM), in concentrations ranging from 5 to 20 at.% [20–23]. The most common PGM addition is Pt. The state-of-the-art Pt-modified aluminide coatings contain approximately 7 at.% Pt [24]. Pt has been suggested to improve the oxidation resistance of aluminide coatings via a number of mechanisms. First, Pt lowers the activity of Al in B2 NiAl. As Pt accumulates near the TGO, due to the selective oxidation of other elements, Al activity gradient is increased, increasing the flux of Al to the TGO/bond coat interface promoting alumina scale formation [25]. Pt additions also improve scale adhesion by inhibiting S segregation to the scale/alloy interface [12,26]. As a result, the thermal cycling lifetimes of Pt-modified aluminide coatings are typically longer than simple nickel aluminide coatings under thermal cycling conditions. As an alternative to Pt, other PGM alloying additions have been investigated [23,27,28]. Of particular interest is Pd, which, while sharing many chemical similarities to Pt, is significantly less expensive on an atom-for-atom basis. However, the degree to which Pd additions improve oxidation resistance is a subject of debate [21,28–31]. The lifetime of Pd-modified coatings are often found to be greater
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Table 1 Nominal bond coat compositions. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
Nominal composition at.% Ni
Al
Cr
Hf
60 56 55 60
34 33 34 23
6 6 6 12
0.3 0.3 0.3 0.3
Pd
Pt
5 5 5
[42]. However, the mechanism by which Hf inhibits rumpling has not been clearly demonstrated. In this work, Hf-modified B2 bond coats with varying Pt, Pd, and Cr additions were thermally cycled to 1150 °C and the oxidation and rumpling behaviors were investigated. The influence of the transformation from B2 to L12 on the system evolution is studied in detail. The performance of these coatings is compared with the stateof-the-art platinum aluminide coatings. 2. Materials and experimental procedure
than unalloyed nickel aluminides, but less than Pt-modified coatings. Also, it is not known how Pd affects the Al activity in B2 NiAl, or S segregation. Even the phase equilibrium in the Ni–Al–Pd system is poorly understood. Rumpling is another phenomenon that can degrade bond coats [32–35]. The low creep resistance of B2 bond coats makes them particularly susceptible to morphological evolution when stressed [32,36–38]. These stresses can develop at operating temperature, as a result of TGO growth stresses, or interdiffusion-induced stresses in the intermetallic. During thermal cycling, stresses also develop due to differences in the coefficient of thermal expansion (CTE) between the TGO, the bond coat, and the superalloy, as well as phase changes occurring during heating and cooling. The relative importance of each of these factors depends upon the details of the thermal history and coating composition. Most B2 coatings have been found to rumple when cycled to maximum temperatures above 1100 °C, with Pt and Al concentrations not strongly affecting the rumpling rate [39–41]. Hf additions, on the other hand, have been observed to decrease the rate of the rumpling
70
a
Ni
Concentration (atomic %)
60 50 40 30 20
Al Cr
10
Pd
0 0
10
20
30
40
50
60
70
Depth (µm) 10
b
Bond Coat
Superalloy 1.5
8 7
Mass Change (mg/cm2)
Sulfur Concentration (appm)
9
Single crystal René N5 (Ni–7.6Co–8Cr–0.9Mo–1.6W–1Re–2.2Ta– 13.7Al–0.05Hf–0.75C–0.02B–0.01Y, at.%) was used as the substrate. Cylindrical disk coupons, 12 mm in diameter and 2 mm in thickness, were coated with four bond coat compositions using ion plasma deposition (IPD) over the entire surface to an average thickness of approximately 50 μm [43]. The coating compositions, measured by X-ray fluorescence (XRF), are presented in Table 1. Three of the coatings, designated 5Pt, 5Pd, and + 5Ni, contained the same concentration of Al, Cr, and Hf with additions of 5 at.% of the metals, Pt, Pd, and Ni, respectively. The fourth coating contained 5 at.% Pd, with less Al but more Cr. Using thermodynamic databases, the Ni and Al contents were adjusted to maintain the B2-NiAl structure. Following deposition, the coatings were heat-treated for 4 h at 1079 °C under vacuum, and then lightly grit blasted (30 psi, 220 grit Al2O3), imparting a surface roughness of approximately 2 μm. The composition depth profile was measured using glow discharge mass spectroscopy, (GDMS), by Evans Analytical Group (Fig. 1). The S concentration profile, also measured by GDMS, is less than 1 atomic ppm (appm) in the bond coat, and less than 5 appm in the superalloy (Fig. 1). The thermal cycle used to oxidize the coatings consisted of a 45 minute dwell period at 1150 °C, with heating and cooling rates of approximately 200 °C/min. Two samples of each bond coat composition, + 5Ni, 5Pd, 5Pt, and 5Pd–12Cr–23Al, were cycled until a net mass loss had occurred, which was defined as failure. A third set of samples underwent 500 cycles. Surface and cross-sectional microstructural analysis was performed using the scanning electron microscope (SEM). The surface morphology of the bond coat specimens was analyzed with an optical surface profilometer, using vertical scanning interferometry, as described elsewhere [44]. Surface height measurements, with a vertical resolution better than 0.1 μm, were scanned over an area of 0.46 × 0.62 mm2 at 10× magnification with a lateral resolution of 1.0 μm, and over an area 0.92× 1.24 mm2, at 5× magnification with a lateral resolution of 1.9 μm. Five measurements were recorded at each magnification. The surface topology was characterized by the root-mean-squared (RMS) roughness, Rq (Eq. (1)), where zi is a measured surface height, z ̅ is the
6 5 4 3 2 1
+5Ni
1.0
5Pd
5Pt
0.5 Pt-Aluminide*[45] 5Pd-12Cr-23Al
0 0
10
20
30
40
50
60
70
Depth (µm)
0.0
0
500
1000
1500
2000
Cycles Fig. 1. The bulk composition profile (a) and sulfur concentration (b) of the as heat-treated 5Pd-12Cr-23Al, measured by GDMS.
Fig. 2. Mass change per unit area of the bond coats thermally cycled at 1150 °C.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
15
Table 2 Number of cycles to maximum mass change and failure, as defined by zero crossing. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
# of cycles Max. mass gain
Failure
267 400 1050 267
975 1050 1775 850
average height, and n is the number of locations where the height was measured in each scan. vffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffiffi u n u1 X 2 Rq ¼ t ðz −z Þ : n i¼1 i
ð1Þ
The RMS roughness can be related to the scale of the surface undulation, by envisioning a sinusoidal surface of the form 2πx 2πy sin zðx; yÞ ¼ A sin L L
ð2Þ
where A is the amplitude, L is the wavelength, and A = 2Rq. 3. Results Bond coat specimens were thermally cycled to a maximum temperature of 1150 °C, and the specific mass change was measured
Fig. 4. Secondary electron SEM image of 5Pt coating after 500 cycles showing cracks in the TGO at the apex of coating ridges.
periodically (Fig. 2). The mass change of grit-blasted, Pt-aluminide coated René N5, reported by Tolpygo et al. [45], is overlaid for comparison. During cyclic oxidation there was an initial period of mass gain, during which the rate of mass gain decreased with cycle number until a maximum was reached, followed by mass loss, during which the rate of mass loss increased with cycle number. The rate of mass change during the initial 100 cycles was similar for all samples tested. Table 2 lists the average number of cycles to maximum weight gain at which the rate of specific mass change became negative, as well as the average number of cycles to failure, as defined by crossing zero.
Fig. 3. Plan view back-scatter SEM micrograph of the coatings after 400 cycles: showing regions of TGO spallation and grain boundary ridges (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt.
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The oxide scale formed on all of the alloys was primarily composed of Al2O3 with fine entrained HfO2 particles, typically 0.1–1 μm in diameter. Surface backscatter SEM images of each bond coat after 400 cycles are presented in Fig. 3, where the dark-contrast phase is Al2O3 and the small bright phases are HfO2. The bright HfO2 particles can be observed on the surface of the 5Pt alloy after 500 cycles (Fig. 4). Cracks in the scale were observed after 500 cycles, with the majority of these cracks residing at the apex of surface undulations (Fig. 4). Oxide spallation was observed emanating from these cracks after 200 cycles in the 5Pd–12Cr–23Al and +5Ni, after 250 cycles in the 5Pd, and after 600 cycles in the 5Pt. Areas of exposed bond coat are evident as the large bright regions in Fig. 5. Areas lacking HfO2 particles, that are roughly the same size as the exposed metal regions, are interpreted to be re-oxidized spalled regions. Higher magnification micrographs of the bond coat surface below spalled oxide (Fig. 5) show imprints of the inwardly growing Al2O3 grains. Oxide pegs, areas of the TGO that grow into the alloy along Hf-rich stringers are evident in the exposed bond coat underneath TGO spalls. Cross-sectional backscattered SEM images after 500 cycles (Fig. 6) confirm that the alloys form an Al2O3 scale with a high density of HfO2 particles near the TGO surface, with occasional larger HfO2 particles throughout the scale. Oxide scales where HfO2 is present at the surface indicated regions that have not spalled and re-oxidized, and the scale thickness in these regions is similar for all alloys, 4–6 μm. The constant scale thickness in non-spalled regions is in agreement with the scale mass change during the initial stages of oxidation.
The frequency of oxide spallation increases concurrently with the decrease in specific mass change. The amount of spalled area is largest for the 5Pd–12Cr–23Al bond coat and smallest for the 5Pt bond coat. Spalled regions are apparent in cross-section (Fig. 7) as regions of bare metal as well Al2O3 absent HfO2, the latter indicative of re-grown oxide. Cross-sectional micrographs of coatings cycled to failure (Fig. 7) show that the oxide scale is still primarily composed of Al2O3 although more of the scale has spalled. Additionally, only a small volume fraction of nickel oxide is observed in the scale. All of the bond coats extensively rumpled during thermal cycling. This is illustrated by the secondary electron imaged surfaces in Figs. 8 and 9. The extent of rumpling can also be observed in the crosssections in Figs. 6 and 7 as well as in the surface contour maps in Fig. 9 that show the 5Pt coating after grit-blasting and after 250 cycles. The RMS roughness as a function of cycle number is presented in Fig. 10. The rate of rumpling was similar for all of the bond coats for the first 100 cycles. The 5Pd–12Cr–23Al coating rumpled the least overall, and little change in the surface roughness was observed after 200 cycles. Rumpling ceased in the 5Pd bond coat after 350 cycles, and the + 5Ni sample experienced no further rumpling after 500 cycles. The 5Pt samples continued rumpling through 1500 cycles, developing with the highest rumpling amplitude of all of the bond coats. The evolution of surface roughening can best be discerned by dividing the lifetime of each sample into three regimes. At the onset of thermal cycling, the roughness of all compositions increases significantly and at a nearly identical rate. Subsequently, the bond coats rumple at different relative rates until a plateau roughness is reached prior to failure. The phase evolution within the bond coats was analyzed by X-ray diffraction (XRD). The initial microstructure of all the coatings consisted solely of B2 NiAl, with the exception of 5Pd–12Cr–23Al, which contained a small fraction of L12 Ni3Al. During cooling in the first thermal cycle, the B2 phase in all samples underwent a martensitic transformation to tetragonal L10 and the monoclinic 7R [46]. During high-temperature exposure, Al is lost from the bond coat due to TGO growth and interdiffusion with the substrate. As a result of Al depletion, the B2 (7R and L10) NiAl phase transforms to the L12 Ni3Al phase. Table 3 lists the time at which L12 is first detected, as well as the point at which the diffraction of L10 is no longer detected. Fig. 11 shows the XRD scans for 5Pd, illustrating the growth of the L12 (111) reflection and the corresponding decrease of the 7R (113) and L10 (200) reflections. 4. Discussion While the initial oxidation and rumpling behaviors of the four alloys were similar, a large variation in the cyclic oxidation lifetime and surface roughness at failure was observed. The discussion below focuses on identifying which compositional factors played a role in the stratification of these phenomena through thermal cycling. 4.1. Cyclic oxidation behavior
Fig. 5. Back-scatter SEM images of +5Ni after 500 cycles. (Top) Spalled regions in which the metallic bond coat surface is visible are outlined in red and re-oxidized regions that do not contain HfO2 are outlined in blue. (Bottom) Higher-magnification images of a spalled area and (inset) HfO2 particles in the TGO, outlined in blue.
The net mass change and oxide morphology during the first 100 cycles are similar for all alloys. Limited spinel formation and no oxide spallation were observed. These results demonstrate that the Al and Cr concentrations in all of the bond coats were sufficient to grow alumina at 1150 °C and that PGM additions are not necessary to prevent the growth of transient oxides, as is the case in bond coats with lower Al and Cr concentrations [6,20,25]. Many alloys exposed to high-temperature thermal cycling fail when their ability to form a protective oxide through selective oxidation is compromised due to compositional change via interdiffusion and/or loss of Al to the oxide [8]. However, alumina growth continued
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
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Fig. 6. Cross-sectional back-scatter SEM images of the (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt bond coats after 500 cycles, showing the rumpling under the TGO scales. Higher-magnification (insets) shows the presence of embedded Hf-rich particles in the TGO.
throughout the thermal cycling test for all four alloys, even as localized spallation and, ultimately, net mass loss occurred. Overlay coatings are not in thermodynamic equilibrium with the substrate, so a flux of elements into the superalloy during thermal cycling is expected. Pt, which decreases the activity of Al in Ni–Al–Pt alloys [47], is expected to decrease the flux of Al into the superalloy by decreasing the Al chemical potential gradient. Pd could be expected to have a similar affect on the Al activity and its interdiffusion behavior. In a parallel set of coatings fabricated by the same approach as described here, the concentration of the Pd in the coating was found to fall to approximately 25% of its original level, while the Pt only fell to about 45% of its original level after 1800 cycles at 1150 °C [Adharapurapu, unpublished research]. Therefore, a lower Al concentration at the TGO interface might be expected in the Pd-containing coatings compared to the Pt-containing coatings [23]. However, because Al2O3 is formed throughout cycling it does
not appear that Al influences the ability of the coatings to reform the protective oxide, within the range of composition employed here. As shown in Fig. 4, the major difference in the response of the four coatings is in the frequency of the spalling. The Pt-containing coating experiences a lower incidence of spalls throughout cycling. Previous high-temperature studies of the mechanical properties of B2 coatings [23] suggest that the differences are dominated by interfacial effects. However, isolating the role of alloying additions on the oxidation of B2 bond coats during thermal cycling is confounded by changes in properties as well as interdiffusion with the substrate. The effect of composition on the Al2O3/bondcoat interface has been the subject of many experimental and theoretical investigations. When the toughness has been directly measured and where the work of adhesion has been calculated through sessile drop tests, it has been found that bulk alloying additions of Cr do not markedly improve
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R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
Fig. 7. Cross-sectional back-scatter SEM images of the bond coats cycled to failure (A) +5Ni, 850 cycles; (B) 5Pd, 1000 cycles; (C) 5Pd-12Cr-23Al, 850 cycles and (D) 5Pt, 1750 cycles.
adhesion [48–50], while Pt additions can increase the work of adhesion by approximately 20% [51]. Smaller improvements can be observed for Y [49], and Al. The effect of composition on interfacial adhesion has been studied more systematically through the use of density-functional theory [52]. These results are in general agreement with the experimental work, and suggest that Cr, Al, Pt, and Hf additions do not markedly increase the work of adhesion between nickel and alumina [19,52]. The metal/Al2O3 work of separation typically falls in the range of 2.5–3.5 J/m 2 and depends on interfacial stoichiometry and the orientational relationship as much as the composition. The influence that alloying Ni with Pd has on the Al2O3/Ni adhesion has not be reported on in the literature, but the work of separation between Pd and Al2O3, is less than Pt, Ni, and Al respectively [52]. Unlike Al, Cr, Pt, Pd, and Hf, sulfur is known to have a strong effect on interfacial toughness [12,15,19,52–55]. There are two primary modes by which sulfur segregation degrades interfacial toughness:
facilitation of the growth of interface voids [5,11], and, the increase in alloy–oxide interfacial energy [53]. Both of these mechanisms can be affected by alloying additions. However, the total S concentration in the bond-coated superalloys is low, less than 1 atomic ppm in the bond coat and 4 ppm in the superalloy (Fig. 1). Nevertheless, with increased exposure time at 1150 °C, S is expected to diffuse toward the TGO due to the large heat of segregation [13,14]. It is therefore important to consider how the alloying additions affect the S segregation. Reactive elements, such as Hf, are known to mitigate the detrimental effect that S has on adherence. HfS2 has a large negative free energy of formation [56] and therefore Hf dissolved in the bond coat will limit the S activity in the alloy. Additionally, Hf segregates to the alloy/alumina interface and because the heat of segregation is larger than that of sulfur, the concentration of sulfur at the interface is limited [19]. However, all of the bond coats investigated contain
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
19
Fig. 8. Plan view secondary electron SEM micrographs of the coatings (A) +5Ni, (B) 5Pd, (C) 5Pd-12Cr-23Al and (D) 5Pt after 400 cycles.
Hf, and a large range in spallation resistance was observed. It is of interest to discuss how Cr, Al, Pt and Pd additions may affect S segregation, and, in turn, adherence. Both theoretical and experimental investigations have shown that Cr promotes the interfacial segregation of S in Ni–Al2O3 systems through a Cr–S co-segregation process. Accordingly, the interface strength of Ni–Al2O3 systems in which sulfur is present is
markedly reduced by Cr additions. Further, Hou has shown that Cr additions can overwhelm the beneficial effects of reactive elements [14,26]. Therefore, the co-segregation of Cr and S to the BC/TGO interface is a plausible explanation for the poor spallation resistance of the 5Pd–12Cr–23Al coating, with respect to the other coatings investigated. Pt additions, widely used in commercial bond coats, have been shown to mitigate the effects of S by reducing the formation of
Fig. 9. Plan view secondary electron SEM micrograph of the 5Pt coating (A) after 0 and (B) 250 thermal cycles and contour map of 5Pt after (C) 0 and (C) 250 cycles showing the topography of the rumpled surface.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
Roughness Rq (μm)
15
5Pt
10 +5Ni 5Pd
5
0
5Pd-12Cr-23Al
0
200
400
600
800
1000
1200
Cycles Fig. 10. RMS roughness of the bond coats thermally cycled at 1150 °C.
interfacial voids and limiting the amount of interfacial S segregation. Accordingly, the 5Pt bond coat showed the greatest spallation resistance. However, the relative importance of Pt and Hf is illustrated when the 5Pt bond coat is compared to the CVD platinum aluminide (Fig. 2) which had a thermal cycling lifetime comparable to the +5Ni sample. Replacing Pt with Pd resulted in a marked decrease in the spallation resistance, but the 5Pd bond coat was still superior to the +5Ni bond coat. The paucity of reported investigations of Pd-modified nickel aluminides system makes it difficult to infer the mechanistic root of the benefit imparted by Pd. Oquad and Monceasu have reported that interfacial voids form when B2 (Ni,Pd)–Al are oxidized at 1100 °C, similar to binary B2 NiAl and in contrast to B2 (Ni,Pt)–Al. However, few voids were observed in the 5Pd bond coat, as is commonly the case for Hf-modified bond coats. It is reasonable to speculate that Pd, like Pt segregates to the metal–Al2O3 interface, inhibiting S segregation and further work to examine this hypothesis would be useful. Finally, it is noted that the concentration of Hf in the coatings was relatively high, and can be considered to be ‘over-doped’ due to the amount of HfO2 present in the TGO. A decrease in the HfO2 concentration from 0.3 to 0.05–0.1 at.% has been shown to increase the cyclic oxidation behavior of NiAl alloys [5] and could improve the oxidation resistance of the coatings investigated in this study.
4.2. Rumpling All of the bond coats in this investigation were found to rumple extensively. The rate at which the surface roughened was similar for all of the alloys during the first 100 cycles. This rumpling behavior is typical of B2 bond coats on René N5 superalloys [36,39]. The primary factors that drive rumpling have been identified as growth strains in the TGO and thermal expansion mismatch strains in the superalloy–bond coat–TGO system [57,58]. These strains should not initially vary significantly among the compositions investigated, given the similarity in the initial oxidation rate and the minor effect that the composition has on the CTE of B2 NiAl bond coats [4], consistent
with the similarity of rumpling rates in all coatings early in the cyclic life. As thermal cycling continues, several phenomena that may affect rumpling are occurring: the B2 to L12 phase transformation, bond coat–superalloy interdiffusion, and oxide spallation. As the Al content in the bond coat decreases, the B2 structure transforms into the L12 phase. This process is expected to affect the rate of rumpling in several ways. First, the B2 phase transforms via a martensitic transformation to L10 during cooling [59]. The volumetric expansion associated with this phase transformation induces stresses that promote rumpling, as illustrated by Balint and Hutchinson [57]. The reduction of the volume fraction of B2 should reduce the amount of strain imposed by the martensitic transformation. The CTE mismatch strain between the bond coat and the superalloy will also change as the phase transformation proceeds. Haynes et al. [60] have reported the average CTE of several Ni–Al alloys including of a B2 alloy, Ni–39Al–5Pt, a L12 alloy, Ni–25Al–Hf, and the two-phase γ/γ′ (L12/FCC) superalloy René N5 (16.0, 15.6, 17.2 ppm K − 1 respectively at 1150°). Therefore, as the B2–L12 transformation proceeds, the CTE misfit strain between the bond coat and the superalloy is expected to increase. If further Al depletion occurs the L12 phase will transform to FCC γ-NiAlCr solid solution. The CTE of γ phase, 18–19 ppm K − 1 [60] is higher the L12, B2 phases and René N5, which will decrease the CTE mismatch strain. Finally, the L12–B2 phase transformation will change the creep behavior of the bond coat. While both B2 and L12 are soft at 1150 °C [61,62], the flow stress of Ni3Al is nearly twice that of NiAl at 800 °C, 180 MPa and 95 MPa respectively [63,64]. Additionally, while the flow stress of B2 decreases monotonically from 800 °C to room temperature, L12 Ni3Al exhibits anomalous behavior with a maximum flow stress around 600 °C [65]. The non-linear temperature dependence of L12 flow stress and the René N5 CTE complicate the analysis of the effect that the B2/L12 volume fraction has on the rumpling rate. The B2 to L12 phase transformation is driven by Al depletion, resulting from Al2O3 growth and interdiffusion with the superalloy, and the volume fraction of the L12 phase increases with time. The rate of L12 growth is greatest in the 5Pd–12Cr–23Al bond coat due to spallation and low initial Al concentration, while the transformation occurs most slowly in the 5Pt bond coat, which exhibits the least oxide spallation and Al back-diffusion [47]. However, the extent of rumpling is also greater for the 5Pt bond coat. The amplitude of the surface undulations may also be affected by oxide spallation, in addition to the thermo-mechanical stress induced deformation of the bond coat. The majority of spallation events result from cracks which form at the apex of surface undulations, as
Intensity (A.U.)
20
Table 3 Number of cycles at which the L12 phase first was detected and after which L10 was no longer detected by X-ray diffraction. Sample designation
+5Ni 5Pd 5Pt 5Pd–12Cr–23Al
# of cycles L12 start
No NiAl
10 10 50 0
400 600 600 400
40
42
44
46
48
2θ
50
52
54
(o)
Fig. 11. X-ray diffraction pattern of 5Pd after 50 and 500 cycles showing the increase in intensity of L12 diffractions peaks corresponding to the increasing volume fraction.
R.W. Jackson et al. / Surface & Coatings Technology 221 (2013) 13–21
illustrated in Fig. 5. These observations are in agreement with the analysis of Hutchinson [66] who calculated the energy release rate for compressed films on curved surfaces, showing that the driving force for delamination is greatest on a convex surface. The selective removal of material at the top of undulations will have the tendency to reduce the overall surface roughness. Accordingly, the order in which spallation was observed in the bond coats, first 5Pd–12Cr– 23Al, followed by + 5Ni, 5Pd, and 5Pt, is the same as the order of observed rumpling rate reduction. The plateau in roughness may therefore be due to a combined effect of spallation and phase transformation. 5. Conclusions A series of Hf-modified B2 bond coats were thermally cycled to a maximum temperature of 1150 °C and the oxidation and rumpling behaviors were characterized. All of the bond coats formed Al2O3 scales. The Pt + Hf containing bond coat was the most spallation resistant and had the longest cyclic oxidation lifetime. The 5Pd–12Cr–23Al coating, which had the lowest Al and highest Cr concentration was the most prone to spallation. The 5Pd and + 5Ni bond coats had similar cyclic oxidation lifetimes as the state-of-the-art, grit-blasted Pt-modified diffusion aluminide bond coats, while the 5Pt lasted nearly twice as many cycles. All of the bond coats investigated were found to rumple. The rate of rumpling decreased with continued thermal cycling. The relative reduction in rumpling rates may be attributed to three factors: (1) the increased creep resistance of the L12 phase, which evolved during high-temperature exposure; (2) the removal of B2 phase which undergoes martensitic phase transformation upon cooling and (3) oxide spallation that, occurring preferentially at convex regions of the rumpled structure, dampened the rumpling amplitude. Acknowledgments The authors are grateful for the support of a NSF-Goali program, DMR Grant DMR-0605700. The authors acknowledge R. DiDomizio with coating processing, V. Dheeradhada for the XRF bond coat composition measurements, and R. Adharapurapu and C. Levi for the helpful discussions. References [1] G.W. Goward, D.H. Boone, Oxid. Met. 3 (1971) 475. [2] A.G. Evans, D.R. Mumm, J.W. Hutchinson, G.H. Meier, F.S. Pettit, Prog. Mater. Sci. 46 (2001) 505. [3] I.T. Spitsberg, D.R. Mumm, A.G. Evans, Mater. Sci. Eng., A 384 (2005) 176. [4] J.A. Haynes, B.A. Pint, Y. Zhang, I.G. Wright, Surf. Coat. Technol. 204 (2009) 816. [5] B.A. Pint, K.L. More, I.G. Wright, Oxid. Met. 59 (2003) 257. [6] C.S. Giggins, F.S. Pettit, J. Electrochem. Soc. 118 (1971) 1782. [7] F.S. Pettit, Trans. Metall. Soc. AIME 239 (1967) 1296. [8] G.H. Meier, Mater. Sci. Eng., A 120 (1989) 1. [9] H.E. Evans, R.C. Lobb, Curr. Sci. 24 (1984) 209. [10] A.G. Evans, G.B. Crumley, R.E. Demaray, Oxid. Met. 20 (1983) 193.
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