Microstructural instability in coated single crystal superalloys

Microstructural instability in coated single crystal superalloys

Journal of Materials Processing Technology 153–154 (2004) 660–665 Microstructural instability in coated single crystal superalloys M. Reid a,∗ , M.J...

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Journal of Materials Processing Technology 153–154 (2004) 660–665

Microstructural instability in coated single crystal superalloys M. Reid a,∗ , M.J. Pomeroy b , J.S. Robinson b b

a Stokes Research Institute, University of Limerick, Limerick, Ireland Materials and Surface Science Institute, University of Limerick, Limerick, Ireland

Abstract Specimens of CMSX-4 and CMSX-10 alloys have been platinum aluminised to give a single phase (Ni,Pt)Al coating and a two phase PtAl2 –(Ni,Pt)Al coating, respectively. Specimens of the coated alloys were exposed in an oxidising environment for 188, 350 and 750 h at a temperature of 1100 ◦ C. Specimens were then sectioned, polished and examined using scanning electron microscopy (SEM), energy dispersive X-ray analysis (EDS) and X-ray diffraction (XRD). Decreases in nickel activity and increases in aluminium activity occurred within the substrate close to the coating/alloy interface. These effects resulted in precipitation of additional ␥ in the alloy, which caused rafting of the ␥ and encapsulation of needle-like sigma phase in ␥ . The development of these needle-like sigma phases beneath the diffusion layer in the CMSX-10 alloy at 1100 ◦ C is significant, such that they extend to a depth of 0.2 mm. In addition, they cause failure at the coating/substrate interface in the case of the CMSX-10 alloy. For the CMSX-4 alloy, equilibration effects remote from the coating cause further precipitation of sigma phases in isolated cells in ␥ regions. The implications of these instabilities on superalloy properties are discussed. © 2004 Elsevier B.V. All rights reserved. Keywords: Superalloys; Diffusion; Platinum aluminide

1. Introduction Since their development in the late 1980s, second generation single crystal superalloys have attained success in both commercial and military aircraft engines [1]. These alloys typically contain 3 wt.% Re. The achievement of microstructural stability in these alloys involves control of topologically closed-packed (TCP) phases that could precipitate in moderate amounts after long exposure times at high temperatures, i.e. 950 ◦ C. TCP phases are not beneficial to superalloys because they weaken the superalloy by providing fracture sites and paths due to their needle-like structure and because they deplete the matrix of strengthening refractory elements [2]. The development of third generation single crystal superalloys proceeded with increases in refractory element content. The development of an early experimental composition, alloy 5A [3] showed that there were new instability issues associated with the additions of higher levels of refractory elements. Not only was the amount of TCP increased over second generation alloys but a new microstructural instability was discovered [4,5]. This detrimental instability occurred under coatings used as remedy against aggressive ∗ Corresponding author. E-mail address: [email protected] (M. Reid).

0924-0136/$ – see front matter © 2004 Elsevier B.V. All rights reserved. doi:10.1016/j.jmatprotec.2004.04.132

environments in which modern gas turbine blades operate. The coatings consist of high concentrations of Al to enhance the oxidation resistance of the superalloys through the formation of Al2 O3 on the surface [2]. The addition of Pt prior to aluminising was found to improve the hot corrosion resistance and scale adhesion [6,7]. The area of instability under the coating was termed either a secondary reaction zone (SRZ) or equilibration zone (equilibration zone) because it occurred beneath the diffusion layer of the coating and involved equilibration of the coating with the substrate. The discovery of the equilibration zone raised new issues in further development of third generation single crystal superalloys. It was proposed that Re was the element most responsible for the formation of the equilibration zone [4]. Thus, the levels of Re and other refractory elements had to be carefully balanced to yield good microstructural stability and mechanical properties, i.e. creep rupture strength. This paper discusses the conditions under which the instability occurs and the effect on properties. 2. Experimental Two alloys were evaluated for their propensity to form an equilibration zone, the second generation CMSX-4 alloy and third generation CMSX-10 alloy, the compositions of which is given in Table 1. The first Re containing al-

M. Reid et al. / Journal of Materials Processing Technology 153–154 (2004) 660–665 Table 1 Nominal compositions (wt.%) of CMSX-4 and CMSX-10 alloys [1] Alloy












60.8 69.6

5.7 5.7

6.7 2.0

9.5 3.0

1.0 0.2

6.4 5.0

3.0 6.0

6.4 8.0

0.5 0.4

0.1 0.03

loys are labelled second generation single crystal superalloys, of which CMSX-4 is a member. The CMSX-4 alloy contains 3 wt.% Re for operation at temperatures typically around 1140 ◦ C. It was designed with the aim of approaching optimal castability, strength, fatigue and creep properties [1]. It consists of approximately 60 vol.% of ␥ cuboidal coherent precipitates with (1 0 0) edges surrounded by a disordered ␥ matrix. More recently, alloy designers have tried to improve the high temperature capability of the single crystal blade alloys by increasing the Re content even further. The CMSX-10 alloy belongs to the third generation of Ni-base superalloys designed for aircraft turbines [1]. The high refractory element content, especially that of Re (6 wt.%), improves the high temperature creep performance of the alloy. The tendency for the precipitation of TCP phases in the third generation alloys is a more important problem than in the second generation alloys since it is difficult to attain the right balance between the alloying elements promoting this instability (Note: CMSX-4 alloy contains 16.4 wt.% refractory elements while the CMSX-10 alloy contains 20.7 wt.%.) Rod shaped specimens of CMSX-4 and -10 alloys approximately 8 mm in diameter and 15 mm long were electroplated with a thin layer of platinum (6–10 ␮m), Pt were diffusion heat treated, subsequently aluminised and heat treated using industrial standard Pt aluminising routes designed to give a single phase (Ni,Pt)Al coating on CMSX-4 alloy and a two phase PtAl2 –(Ni,Pt)Al coating on CMSX-10 alloy, respectively, by SIFCO Ireland∗ . For isothermal exposure, samples were placed in a muffle furnace for 188, 375 and 750 h at 1100 ◦ C. Specimens were placed in cylindrical quartz cru-


cibles in order to collect spalled oxide and prevent any mechanical damage occurring. A JOEL JSM-840 scanning electron microscope equipped with a Princeton Gamma-Tech (PGT) energy analysis dispersive X-ray (EDS) system was used to obtain high magnification electron images and carry out chemical analysis of different phases present. The samples were observed using backscattered electrons at 20 kV. The EDS spectra were collected at 20 kV, a beam current of 0.26 nA and a working distance of 39 mm, which was reproducible to a high degree of accuracy by focusing standards and specimens using an integrated optical microscope. The spectrum were processed using PGT software, using standard reference spectra obtained from pure elements under the same conditions. For Re a standard reference spectra was obtained using CMSX-10 alloy and employing analysis results obtained from a fully calibrated X-ray fluorescence (XRF) instrument. The spectrum lines that were used are Ni–K, Cr–K, Co–K, Al–K, Ti–K, W–M, Re–L, Ta–L, Mo–L, Hf–L, and Pt–L. A Philips X’Pert XRD was used in conjunction with the EDS analysis results and backscattered electron imaging contrast to identify phases present. Ni filtered Cu K␣ radiation was used and diffractograms were obtained over a scanning range 20–90◦ 2θ with a scan speed of 0.005◦ 2θ/s. The X-ray diffraction experiments were calculated for an X-ray wavelength of λ = 1.54056 nm and the angle of incidence of the beam on the surface varied from θ =10–45◦ . Under such conditions the structural information obtained from X-ray diffraction patterns corresponds to a surface layer approximately 16 ␮m thick.

3. Results Fig. 1 displays typical equilibration zone regions found beneath platinum aluminide coatings on CMSX-4 and CMSX-10 alloys. For the CMSX-4 alloy the instability was observed as isolated cells of bright needle phases in a dark matrix (Fig. 1(a)) at the substrate coating inter-

Fig. 1. Backscattered electron images showing the equilibration zone after 375 h at 1100 ◦ C on: (a) CMSX-4 alloy and (b) CMSX-10 alloy.


M. Reid et al. / Journal of Materials Processing Technology 153–154 (2004) 660–665

Fig. 2. Backscattered electron image showing the equilibration zone after 750 h at 1100 ◦ C on CMSX-4 alloy.

face after 375 h isothermal exposure at 1100 ◦ C. Similar needles were observed in the CMSX-10 alloy after 375 h isothermal exposure, however, the needles precipitated in a continuous dark matrix (Fig. 1(b)). The bright needle phase in the backscattered electron images (Fig. 1) indicates a high refractory element content. Elemental mapping and quantitative EDS of the needles revealed the phases to be primarily Re-rich and also to contain Cr and W. The phase composition would indicate either P or ␴-phase, of presumably a similar composition and structure to the phases observed by Rae and Reed [8]. Numerous needle-like phases had precipitated in the centre of the alloys between 188 and 375 h exposure at 1100 ◦ C with the percentage and size of such phases increasing after 750 h isothermal exposure. The precipitates in the centre of the alloy are more likely a result of alloy equilibration than any coating/substrate equilibration effect. The precipitates of particular interest in this work are those located at the alloy/coating interface. Fig. 2 shows the equilibration zone/substrate interface of the CMSX-4 alloy after 750 h isothermal exposure, showing the needles precipitating out within ␥ regions. The equilibration zone in which the needles have precipitated is of uniform dark contrast in comparison to ␥ + ␥ matrix in the substrate, indicating that the phases are precipitating in the ␥ -phase. In an attempt to further identify the needles and the matrix in which the needles precipitated, the coating surface on the CMSX-10 alloy was ground down and polished until the matrix was clearly visible, thus making subsequent XRD analysis of the needles and matrix possible. The equilibration zone in which the needles precipitated was confirmed to be ␥ and the bright needle-like phases ␴-phase. The ␥ + ␥ matrix no longer comprises cuboidal ␥ precipitates in a disordered ␥ matrix. The ␥ precipitates have coalesced to form a lath-like structure, commonly referred to as rafting [9–12]. As a result of the extensive needle precipitation, cracking was observed at the coating/equilibration zone interface

Fig. 3. Backscattered electron image showing a crack running along the coating equilibration zone interface after 375 h at 1100 ◦ C on the CMSX-10 alloy.

on the CMSX-10 alloy. The cracking initially was observed after 375 h at 1100 ◦ C, Fig. 3 shows the progress of the cracks along the coating equilibration zone interface. And after 750 h at 1100 ◦ C a similar crack developed approximately 1.6 mm in length and arrested approximately 20 ␮m from the coating surface. No coating delamination was observed so it is possible the cracking occurred during sample removal from the furnace and subsequent thermal expansion mismatch between the ␴-phase and matrix.

4. Discussion 4.1. Nucleation and growth A quantitative treatment of the nucleation phenomenon is proposed in this section to support the microstructural observations. A difficult aspect in studying the nucleation of these constituents is there appears to be a large barrier to their formation. Therefore, predicting formation as a function of time and temperature is difficult and observing the earlier stages of nucleation is without a solution at the present. The total free energy change for nucleation of a nucleus of general shape [13]: G = V(G␥/␥ ) + ηαV 2/3


where η is the shape factor, α the surface free energy between the matrix and the precipitate interface, and V a volume term. Nucleation and growth of solid inclusions in a matrix involves elastic strain energy, Gε . The elastic strain energy reduces the effective driving force of the reaction, G␥/␥ . Where G␥/␥ corresponds to Gibbs free energy change per unit volume. A number of factors control the nucleation of the equilibration zone; surface energy, strain energy, number of heterogeneous sites, and supersaturation. The strain energy can be introduced by surface preparation

M. Reid et al. / Journal of Materials Processing Technology 153–154 (2004) 660–665

Fig. 4. Relationship between the square root of time, temperature, and equilibration zone thickness on CMSX-10 alloy at 1100 ◦ C (␴-phase).

prior to coating, in this case prior to electroplating of the platinum layer. Common surface preparation techniques include shot peen, grit blast, and grinding. Supersaturation can occur either by external or internal chemistry imbalance, i.e. Ni, Al and Pt. The increased Al content due to the inward Al diffusion and outward Ni diffusion decreases the solubility of Re, W, Cr, etc. in the ␥ -phase [14]. This increases the driving force for the precipitation of Re-rich, ␴-phase. The growth of the needle phases at the substrate/coating interface on the CMSX-10 alloy was measured and Fig. 4 shows the data plotted as a function of the square root of time. The linear dependence shows that diffusion is controlling the rate of growth. The countercurrent diffusion of Ni and Al must therefore explain the growth of the needles, which are surrounded by ␥ . Because of these diffusion effects, the activity of substitutional solid solution additives (Re, Cr, W) increase to levels where they precipitate out as intermetallics, which is consistent with observations by Blavette et al. [14], where with the increased Al content the solubility of Re, W and Cr decreases in the ␥ -phase. Cuboids of ␥ -Ni3 Al also become linked as Al diffuses in the ␥ channels between them forming rafts of ␥ and ␥ , as shown is Fig. 2. With the Ni diffusion outwards from the ␥-rich channels and Al diffusion inwards enabling ␥ -Ni3 Al formation. Both effects favour ␥ formation and precipitation of ␴-phase since the solubilities of Re, W and Cr decreases in the ␥ -phase [14–16] and local decreases in nickel concentrations induce precipitation of ␴-phase in ␥ channels. The CMSX-10 alloy is clearly more strongly affected by coating alloy equilibration than the CMSX-4 alloy. In the CMSX-4 alloy the instability was observed as isolated cells of needle phases in a ␥ matrix (Fig. 1(a)) at the substrate coating interface after isothermal exposure at 1100 ◦ C. Similar needles were observed in the CMSX-10 alloy after isothermal exposure at 1100 ◦ C, however, the needles precipitated in a continuous ␥ matrix (Fig. 1(b)). The different instability at the coating alloy interface can be attributed to the difference in alloy chemistries. The higher Ta level in the CMSX-10 alloy will have a stabilising effect on the ␥


matrix since, Ta partitions preferentially to ␥ substituting for aluminium in Ni3 Al structure [17]. The Co content in the CMSX-10 alloy is less than in the CMSX-4 alloy. The CMSX-10 alloys Co content was limited so to reduce the tendency to form TCP phases [18]. However, Waltson et al. [19] recommended higher Co content in order to improve phase stability. A recent investigation by Rae and Reed [8] has shown that substituting Co for Ni produces an increase in the ␥ lattice parameter decreasing the rate of nucleation of ␴-phase, this clearly indicates the importance of Co in increasing the stability of single crystal alloys. Chen and Little [20] attributed the precipitation of needles in ␥ matrix to platinum diffusion inwards, concluding that the size of the equilibration zone was consistent with the extent of platinum diffusion. However, the work carried out in this study indicates that Ni, Al diffusion and alloy chemistries are the primary cause of needle precipitation in a ␥ matrix. 4.2. Implications of equilibration zone formation The presence of such a large quantity of ␴-phase is serious worry in superalloys. The precipitation of needle-like phases will eventually lead to coating failure as a result of a combination of negative effects: (a) Withdrawal of the solid solution strengthening elements (i.e. Re, W, Cr) from the ␥ matrix [2,8,21,22]. (b) Change in alloy chemical segregation could alter lattice mismatch [21]. (c) The disruption of the matrix also causes the emergence of a ␥/␥ rafted microstructure [9–12]. (d) Plate-like structures provide initiation sites for fracture and subsequently a fracture path following a zig-zag along plates [2,22]. In almost all cases in Ni alloys TCP phases occur in a plate-like morphology, which appears as needles in a single phase microstructure [2]. This plate-like structure is significant since it affects mechanical properties negatively. The ␴-phase is generally believed to be detrimental in both crack initiation and propagation [8]. Fig. 3 shows cracking at the equilibration zone/coating interface clearly demonstrating fracture as a result of the brittle ␴-phase. Clearly, the development of needle-like structures in the microstructure will give rise to local stress enhancement and since the thermal expansion coefficients of the CMSX-10 alloy and intermetallics coatings are different, thermally induced stresses arising on cooling could be concentrated at the sharp changes in section with the result that crack nucleation and growth might be expected to occur. The formation of the TCP phases will also cause withdrawal of the solid solution strengthening elements, which in turn will cause a decrease of strength in the matrix [8,21,22]. Depending on conditions of stress and temperature, rafting may have either a detrimental or beneficial effect on


M. Reid et al. / Journal of Materials Processing Technology 153–154 (2004) 660–665

creep behaviour [23–25]. These conditions can broadly be classified as those under which the precipitates are sheared and under which they are not sheared, respectively [23]. Under conditions of constant load and temperature where the rupture life is sufficiently long to allow rafting, the morphological change appears to increase creep resistance [25]. This is necessary because during constant load and temperature testing if sufficient time exists for rafting to occur the stress must be relatively low. This reduces the possibility of precipitate shear, if, however, a specimen is crept at high temperature and low stress, until fully rafted, and then the temperature reduced and load increased to conditions where the precipitates are sheared. The lower temperature resistance is usually reduced compared with the cuboidal structure [23]. In typical a gas turbine environment it is expected that rafting would have a deleterious effect on mechanical properties of the alloy. The presence of the equilibration zone is a serious issue for all newer generation alloys with high percentage refractory elements which require platinum aluminide coating applied as either a bond coat or protective coating. The easiest method to reduce or eliminate the occurrence of these constituents is to change the alloy composition. Lowering the level of refractory element content in the alloy, especially Re will eventually eliminate the formation of such undesirable phases. Unfortunately this is not a viable option, because these alloys are designed for high creep rupture strength and reducing the refractory element content will have a negative effect on the strength of the alloy in this respect. A better understanding of the driving force for equilibration zone development and the effect of the alloy composition on its rate of formation is required to develop an alloy that balances strength and microstructural stability.

5. Conclusions The precipitation of ␴-phase is due to decreases in Ni activity caused by diffusion from the substrate into the coating. Ni activity is further decreased by Al diffusion into the substrate. The decreases in Ni activity and increases in Al activity result in precipitation of additional ␥ , which results in rafting and encapsulation of needle phases by ␥ . The formation of the sigma phase and rafting are of significance since they are detrimental to the mechanical properties of the alloy. The formation of needle-like sigma phase causes failure at the coating/substrate interface for the CMSX-10 alloy.

Acknowledgements The authors are grateful to SIFCO Ireland Ltd. for coating the alloys used in this study. Materials Surface Science Institute (MSSI) supported the present investigation.

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