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Mechanism of failure in a free-standing Pt–aluminide bond coat during tensile testing at room temperature
Materials Science and Engineering A 527 (2010) 842–848
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Mechanism of failure in a free-standing Pt–aluminide bond coat during tensile testing at room temperature Md. Zafir Alam a,∗ , B. Srivathsa a , S.V. Kamat a , V. Jayaram b , N. Hazari a , D.K. Das a a b
Defence Metallurgical Research Laboratory, Hyderabad 500 058, India Indian Institute of Science, Bangalore 560 012, India
a r t i c l e
i n f o
Article history: Received 29 June 2009 Received in revised form 30 August 2009 Accepted 8 September 2009
Keywords: Bond coat Pt–aluminide Microtensile test Failure mechanism Crack nucleation and propagation Fracture toughness
a b s t r a c t The room temperature (RT) tensile behaviour of a free-standing high activity Pt–aluminide bond coat has been evaluated by microtensile testing technique. The coating had a typical three-layer microstructure. The stress–strain plot for the free-standing coating was linear, indicating the coating to be brittle at RT. Different fracture features were observed across the coating layers, namely quasi-cleavage in the outer layer and inner interdiffusion zone, and cleavage in the intermediate layer. By employing interrupted tensile test and observing the cross-sectional microstructure of the tested specimens, it was determined that failure of the microtensile samples occurred by the initiation of a single crack in the intermediate layer of the coating and its subsequent inside-out propagation. Such a mechanism of failure has been explained in terms of the fracture features observed across the sample thickness. This mechanism of failure is consistent with fracture toughness values of the individual coating layers. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Nickel based superalloy components such as blades and vanes that operate in the hot sections of gas turbine engines are usually applied with Pt–aluminide (PtAl) bond coats for enhancing their high temperature oxidation resistance [1,2]. These coatings are formed on the superalloy first by depositing a layer of Pt and then subjecting the plated superalloy to a vacuum diffusion treatment. Subsequently, the superalloy is given an aluminizing treatment to form the coating. Depending on the aluminizing technique used, these coatings are classified as inward grown or outward grown [1,3,4]. These two varieties are also called as high activity and low activity coatings, respectively. The bond coats, which are typically 50–100 m thick, are based on the B2-NiAl intermetallic system. Therefore, they are inherently brittle and have a high brittle-to-ductile transition-temperature (BDTT) above 600 ◦ C [5–11]. PtAl coatings have been reported to degrade the mechanical properties, especially the tensile ductility, of the substrate superalloy at temperatures below the BDTT [5–7].
∗ Corresponding author at: Surface Engineering Group (Coatings), Defence Metallurgical Research Laboratory, P.O.-Kanchanbagh, Hyderabad 500 058, India. Tel.: +91 40 24586634; fax: +91 40 24340683. E-mail address: zafir [email protected] (Md. Z. Alam). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.09.017
In most of the early literature, the effect of bond coats on the mechanical properties (including tensile) of the superalloys had been studied by conducting tests on the coated superalloy samples [5–8]. In recent years, however, it has been realized that it is important to understand the mechanical behaviour of the stand-alone coatings in order to gain better insight into their effect on the properties of the superalloys. Hemker and co-workers [9,10] have reported the tensile properties of outward grown (low activity) free-standing PtAl coating between room temperature and 1000 ◦ C. They have adopted microtensile testing technique to evaluate the tensile properties of these PtAl coatings. Their studies have established several important aspects of low activity PtAl coatings such as their BDTT being about 600 ◦ C. Above this temperature, ductility of the coating has been found to increase considerably while the strength drops very rapidly. Similar values of BDTT of PtAl coatings have also been evaluated by Eskner and Sandström [11] by using a different test method, i.e. miniaturized disk bending technique. Despite the above-mentioned reports on the tensile behaviour of free-standing PtAl coatings, to the best of our knowledge, there has not been any study correlating the microstructure of PtAl coatings with their tensile failure behaviour. In the present investigation, such a correlation has been made in case of a stand-alone high activity (inward grown) PtAl coating at RT. Based on this correlation, a mechanism of failure in the coating has been proposed. Microtensile test method has been adopted to determine the tensile properties of the coating.
Md. Z. Alam et al. / Materials Science and Engineering A 527 (2010) 842–848
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Fig. 1. Schematic of a microtensile sample. All dimensions are in mm.
2. Experimental details 2.1. Sample design Microtensile samples of the geometry as shown in Fig. 1, were used in the present study. The thickness of the sample, which was equal to the coating thickness was approximately 100 m. For the samples of the above geometry, conventional gripping of the specimen between two flat plattens invariably led to the breakage of the samples at the fillet region. This was caused due to the inherent brittleness of the coating. To overcome this problem, the samples were held in slotted grips for testing as shown in Fig. 2. In order to ensure the failure of the samples within the gage length, use of appropriate holding length, HL, and fillet radius, R, (Fig. 1) was found to be critical. To fix these two dimensions, a finite element method (FEM)-based elastic analysis was carried out. Using the appropriate boundary conditions imposed by the slotted grips, a combination of HL and R was determined such that, upon application of tensile loading, a uniform maximum stress was generated within the gage length while the elastic stress concentration at the fillet remained at a minimum value. It was determined that a combination of R = 0.5 mm and HL = 1 mm, which corresponded to an elastic stress concentration at the fillet of 1.02, would lead to failure of the microsamples within the gage length. Based on this analysis, samples with gage length (GL) of 2 mm, gage width (GW) of 0.5 mm, overall length of 8 mm, HL of 1 mm and R of 0.5 mm were prepared. Initial tensile tests indicated that the samples with the above dimensions failed within the gage length on most occasions. Therefore, these sample dimensions were used for evaluation of coating properties.
Fig. 2. A free-standing PtAl microtensile sample placed in the slotted grips while testing.
Fig. 3. Method adopted for the fabrication of microtensile specimens: (a) PtAl coated 0.5 mm thick superalloy strip, (b) wire EDM machining of samples from the coated strip and (c) a typical cut-out microtensile sample.
2.2. Sample preparation, tensile testing and microstructural characterization Directionally solidified (DS) rods of Ni-based superalloy CM247LC were used as the substrate material for developing the PtAl bond coats. The nominal composition (in wt.%) of the above alloy is 9.2 Co–8.1 Cr–9.5 W–5.6 Al–3.2 Ta–1.5 Hf–0.7 Ti–0.015 Zr–0.5 Mo–0.15 B–0.07 C–balance Ni. The coating was formed on 0.5 mm thick strips sliced from the above DS rods. For coating formation, a 5 m thick Pt layer was first electrodeposited on these strips. The plated strips were then subjected to a diffusion heat treatment in vacuum, following which they were given a pack aluminization treatment. A two-step high activity pack aluminizing process [1,2,12,13] was used in this study. Microtensile samples were cut from the coated strips by electro discharge machining (EDM) as shown in Fig. 3(b). The gage length of the microsamples was always parallel to the longitudinal axis of the strips, i.e. along 0 0 1 solidification direction (see Fig. 3(a)). The above cut-out samples (Fig. 3(c)) were then precisionpolished from one side using polishing papers of 600, 1000, 1500 and 2500 grades, respectively. Final finishing was carried out on polishing films containing 1 and 0.5 m diamond particles. The unpolished surface of the coating was comparatively much rougher having an average roughness (Ra ) of 1.3 m as compared to the 0.5 m finish of the polished surface (see Fig. 1). To prepare free-standing coating samples, polishing was carried out till the substrate was completely removed and the sample thickness became equal to the thickness of bond coat, which was approximately 100 m. A Walter–Bai Ag microtensile testing machine with 500 N loading capacity was employed for testing these freestanding coating samples at RT. Five coating microsamples were tested at an initial strain rate of 4.1 × 10−3 s−1 and a constant crosshead velocity of 0.5 mm/min. A pre-stress of 10 MPa was applied for proper gripping of the microsamples within the slotted grips (Fig. 2). Microstructural and fractography observations were made in a Quanta 400 scanning electron microscope (SEM) operating at 20 kV. A Cameca SX-100 electron probe micro analyzer (EPMA) operating at 20 kV was used for chemical analysis of the coating. Identification of the phases present in the coating was carried out by X-ray diffraction (XRD) method using a Philips PW-3020 X-ray diffractometer. Indentation studies were carried out using a Leica VHMT Auto microhardness tester and a CSM instrumented nanoindenter.
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Fig. 4. Microstructure of the as-formed coating showing the individual layers. Some of the voids present in the outer layer are indicated by arrows. Ni plating was given on the sample for edge-protection during metallographic polishing. 1 and 2 correspond to locations whose compositions have been indicated in the ternary phase diagram (Fig. 5).
3. Results
Fig. 6. Variation in Ni, Pt and Al concentrations across the coating thickness.
3.1. Coating microstructure The PtAl coating, as shown in Fig. 4, revealed a three-layer microstructure, typically obtained in a high activity aluminizing process [1,2,12,13]. The matrix phase in all the layers was B2NiAl having Pt in solid solution, i.e. B2-(Ni,Pt)Al. The outer layer additionally contained PtAl2 phase and numerous fine precipitates. The intermediate layer had a single-phase B2-(Ni,Pt)Al structure and contained comparatively much fewer precipitates. The B2(Ni,Pt)Al + PtAl2 structure of the outer layer and the B2-(Ni,Pt)Al structure of the intermediate layer (indicated by points 1 and 2, respectively, in Fig. 4) are consistent with the corresponding compositions of these layers, which can be confirmed based on the ternary phase diagram of Ni–Pt–Al reported by Gleeson et al. [14]. As shown in Fig. 5, the composition of a typical location in the outer layer (point 1 in Fig. 4) corresponds to the B2-(Ni,Pt)Al + PtAl2 phase field. Similarly, the composition of the intermediate layer (point 2 in Fig. 4) lies in the single-phase field of B2-(Ni,Pt)Al. The inner layer of the coating was the interdiffusion zone (IDZ), typically found in most diffusion aluminide coatings [1,2]. This layer is formed as a
result of the loss of Ni from the superalloy substrate required for coating formation [1,2,12,13]. Such loss of Ni leads to precipitation of numerous complex precipitates [1,2,12,13] as seen in the IDZ of the present coating (Fig. 4). The thicknesses of the abovementioned three layers were approximately 60, 25 and 20 m, respectively. The fine precipitates found in the outer layer of the coating were rich in Ni and W [13] and identified as orthorhombic NiW by XRD method [15]. A few irregular shaped Kirkendall voids were also observed in the top regions of the outer coating layer (Fig. 4). The coating did not contain any crack in as-coated condition. Being a diffusion coating, the present PtAl coating contained almost all the substrate alloying elements such as Cr, W, Ta, Hf, Mo and Co. The concentration profiles for Ni, Al and Pt, as shown in Fig. 6, revealed the graded composition of the coating. The Ni concentration increased continuously from about 25 at.% to nearly 50 at.% over the outer and the intermediate layers. Its concentration remained in the range 40–50 at.% in the IDZ. The Al concentration remained in the narrow range of 45–40 at.% over the outer and intermediate layers and fell to almost 20 at.% at the IDZ/substrate interface. The Pt concentration varied fairly sharply across the coating thickness with the outer layer, intermediate layer and the IDZ containing a maximum of about 10, 3 and 0.1 at.% Pt, respectively. 3.2. Tensile behaviour
Fig. 5. Isothermal section of Ni–Pt–Al ternary system at 1100 ◦ C [14]. Composition of locations 1 and 2 correspond to the compositions of regions marked 1 and 2 in Fig. 4.
The stress–strain plot for the free-standing PtAl coating, as shown in Fig. 7, was fairly linear and did not show any appreciable plastic yielding. The total fracture strain was determined to be about 0.2% and the plastic strain was negligibly small. Such low plasticity for the PtAl coating at RT is expected considering their high BDTT values, which are usually above 600 ◦ C [9,10]. The elastic modulus and the fracture stress were determined to be 100 ± 1 GPa and 215 ± 40 MPa, respectively. The linear stress–strain relation observed in the present study is similar to that reported for low activity PtAl coatings by Pan et al. [9]. The typical fracture surface of a failed coating microsample is presented in Fig. 8(a). As evident from this figure, the three coating layers showed distinctly different fracture features. The outer layer and the IDZ showed quasi-cleavage features (Fig. 8(b) and (d), respectively), while the intermediate layer exhibited cleavage features with river-bed patterns (see Fig. 8(c)).
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Fig. 7. Stress–strain plot for the free-standing PtAl coating.
3.3. Crack initiation and propagation After the completion of the tensile test, the failed pieces of the tensile sample were carefully examined for the presence of any crack on both as-coated and polished side (see Fig. 1) in SEM. Not even a single crack in the transverse direction, i.e. normal to the loading direction, could be found over the gage length on both the surfaces of the sample, as typically shown in Fig. 9(a) and (b). To confirm this fact further, metallographically polished longitudinal sections of the failed pieces (see Fig. 1) were also examined. No cracks (perpendicular to the coating thickness) in these sections could be located over the entire gage length, as evident from Fig. 9(c). Further, no crack could also be located on the edges on
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both polished and unpolished surfaces. This is an indication that cracking perpendicular to the loading direction did not occur on the thickness side of the microtensile sample. These observations clearly suggest that the failure of the sample during testing has most likely occurred by the generation and propagation of the first crack in the sample. To investigate this aspect further, interrupted tensile test on four different microtensile samples was conducted wherein the samples were loaded to stress levels of 50, 100, 150 and 180 MPa, respectively, and subsequently unloaded. Care was taken to prevent any breakage of the samples during removal from the slotted grips (Fig. 2). Subsequently, the longitudinal section of the samples was examined for the presence of any crack. The samples subjected to stress levels below 150 MPa did not exhibit any cracking through out the gage length. However, for samples subjected to stress levels of 150 MPa and beyond, only a single crack over the entire gage length was located as shown in Fig. 10(a) and (b). For the sample loaded to 150 MPa, the crack could be seen initiating in the intermediate layer of the coating (Fig. 10(a)). However, the sample loaded to a higher stress of 180 MPa revealed a crack that had propagated across the outer and the intermediate layers, and partially across IDZ (Fig. 10(b)). Even in this sample, the higher crack width in the intermediate layer of the coating suggests that the crack has initiated in this layer and propagated subsequently to the other two layers. The above observations can be summarized as follows. The failure of the stand-alone coating microtensile sample has occurred by the initiation and subsequent growth of the first crack in the sample. The crack has initiated in the precipitate-lean intermediate layer of the coating and subsequently propagated inside-out leading to the sample failure. In order to substantiate the above observation, the fracture toughness across the individual layers of the coating was estimated. For this estimation, a series of indentations were made in each of the coating layers using a Vickers indenter at various loads. Subsequently, they were observed in
Fig. 8. (a) Fracture surface observed on the failed microtensile sample showing the variation of fracture features across the coating thickness. Magnified view of the fracture surface in the outer layer (OL), intermediate layer (IL) and the IDZ are shown in (b), (c) and (d), respectively.
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Fig. 9. Micrographs depicting the various surfaces of the failed tensile sample: (a) as-coated top side, (b) polished side and, (c) longitudinal section of the gage. No crack was located on any of these regions of the sample.
Fig. 10. Longitudinal section of the coating microsamples subjected to interrupted tensile test in which the loading was stopped at (a) 150 MPa and (b) 180 MPa.
SEM to examine the presence of any cracking around the indentations. It was found that between 2 and 3 N loads, lateral cracks originated around the indentations. The minimum load at which cracking around the indentation was observed in any given layer was considered as the critical load for that layer. Since lateral cracking was observed around the indentations, the fracture toughness value was calculated from the relation proposed by Marshall et al. [16], as given in Eq. (1):
4 KIND =
Po A2o 2/3 (cot ) ıo
H4 E
(1)
In the above equation, KIND is the indentation fracture toughness, Po the threshold load for lateral cracking, Ao and ıo the dimensionless constants, the half included angle of the indenter, H the hardness of the material in GPa and E the indentation modulus in GPa. Using the appropriate values of the above-mentioned terms in the equation for each of the three coating layers, the fracture toughness for the individual layers was calculated, as presented in Table 1. It is evident that the fracture toughness of the intermediate layer at 1.15 MPa m1/2 is the lowest among the three coating layers, which supports the earlier observation of the crack initiating in the intermediate layer during tensile test.
4. Discussion As mentioned earlier, the general linear tensile behaviour of the present high activity PtAl coating (Fig. 7) was similar to that reported for low activity PtAl coating [9]. Further, the elastic modulus for the present coating at 100 ± 1 GPa was also similar to the value of 117 GPa reported by Pan et al. for the low activity coating [9]. However, the fracture strength of 215 MPa for the present coating was much higher than the value of 100 MPa reported for the low activity coatings [9]. The difference in the fracture strength values in the two studies can be ascribed to the difference in the stoichiometry and the associated defect strengthening
Table 1 Indentation fracture toughness, KIND of various layers of the Pt–aluminide coating. The data shown has been averaged over ten readings in each layer. The values of dimensionless constants Ao and ıo are 3/4 and 1200, respectively. The half included angle for the Vickers indenter was 68◦ . The values of H, E and KIND are the average of ten measurements.
Outer layer Intermediate layer IDZ
Po (N)
H (GPa)
E (GPa)
KIND (MPa m1/2 )
2 3 2
8.39 ± 0.64 4.81 ± 0.32 7.33 ± 0.17
302 ± 71 244 ± 28 411 ± 31
1.71 ± 0.13 1.15 ± 0.09 1.36 ± 0.02
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Fig. 11. Variation in the stoichiometry of the B2-(Ni,Pt)Al phase present in high activity and low activity coatings. The dashed line indicates stoichiometric (Ni,Pt)50 Al50 .
of the B2-NiAl phase present in the matrix in both the coatings. It is well known that the stoichiometric structure, i.e. Ni50 Al50 structure, is least defective in case of binary B2-NiAl phase [17,18]. On Al-rich side of the stoichiometry, i.e. for hyperstoichiometric structures, vacancy defects develop because of the unfilled Ni-sites. For hypostoichiometric (Ni-rich) structures, however, substitutional defects develop because of excess Ni atoms occupying Al lattice sites [17,18]. It has been widely reported that that the presence of such defects causes some degree of strengthening in bulk B2-NiAl. Further, vacancy defects in hyperstoichiometric B2-NiAl cause much greater strengthening as compared to the corresponding substitutional defects in hypostoichiometric B2 phase [17,18]. High activity coatings are generally much higher in Al content as compared to low activity coatings. Therefore, the B2 phase in the high activity coatings is expected to be largely hyperstoichiometric in structure. To examine this aspect in the present high activity coating, it was assumed that (Ni,Pt)50 Al50 is the stoichiometric composition for the B2 phase, since it is known that Pt occupies Ni-sites in the B2 lattice [19]. Assuming that the effect of site occupancy of other elements in the B2 phase can be neglected, the off-stoichiometric composition can be written as (Ni,Pt)50+d Al50−d , where d is a parameter (in at.%) representing the deviation from stoichiometry. The value of d is zero for stoichiometric composition, positive for hypostoichiometric composition and negative for hyperstoichiometric compositions. The value of d can be determined by equating (50 + d)/(50 − d) to the atom ratios of (Ni + Pt) to Al, as determined from the EPMA composition profiles (see Fig. 6). The variation of d corresponding to the single-phase B2-(Ni,Pt)Al across the thickness of the present coating has been plotted in Fig. 11. The B2-(Ni,Pt)Al + PtAl2 two phase region (approximately 25 m in thickness) of the outer layer has not been considered in the above plot (see Fig. 11) since it was impossible to accurately determine the composition of the B2 phase in this region using the EPMA. For comparison, the corresponding plot for a typical low activity coating (MDC150L) has also been given in the same figure. The plot for the MDC150L coating was carried out using the data reported by Das et al. [20]. It is clear from Fig. 11, that the B2 composition up to a depth of about 60 m in the present coating is hyperstoichiometric, whereas it is completely hypostoichiometric in case of MDC150L coating. The comparative higher fracture strength of 215 MPa for the present PtAl coating is most likely because of its hyperstoichiometric composition and the cor-
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responding high vacancy defects as compared to the low activity coating. However, since modulus is not affected significantly by the deviation from the stoichiometry [17,18], the modulus of the present coating was found to be similar. As previously mentioned, there was no crack present in the coating in the as-coated condition, although some voids in the outer layer could be observed (Fig. 4). Therefore, it can be concluded that the crack causing the failure of the microtensile sample has initiated during the tensile test. The microstructural evidence (Fig. 10(a) and (b)) as well as the fracture toughness data of the individual coating layers (Table 1) indicate that the above crack has initiated in the intermediate layer. The predominantly cleavage fracture in the intermediate layer of the coating (Fig. 8(c)) further supports the above conclusion. Thus, all the evidences, namely the origin of the crack as seen in the coating cross-section (Fig. 10(a)), fracture toughness values and the fracture surface features, together clearly suggest that the failure of the microtensile sample has occurred by initiation of a single crack in the intermediate layer which subsequently propagated inside-out (Fig. 10(b)). It is interesting that despite the presence of stress raisers such as voids close to the surface of the sample (Fig. 4) and also appreciably large roughness of 1.3 m on the unpolished side of the microtensile sample, the cracking initiated in the sub-surface intermediate layer (Fig. 10(a) and (b)). The first crack in the sample leading to failure clearly suggests that once a crack is initiated, its propagation is extremely easy in the brittle B2-(Ni,Pt)Al matrix of the coating. Thus, there is no apparent driving force for the generation of any additional crack in the sample during testing. From the above discussion, it appears that the overall strength of the coating would be primarily dependent on the strength of the intermediate layer. Therefore, strengthening of this layer would increase the overall strength of the coating. The strength of the intermediate layer can be enhanced through refining the grain size of the B2(Ni,Pt)Al phase constituting this layer. Increased Pt solid solution strengthening of the above phase may also be possible by enhancing the Pt content of the coating. In fact, our preliminary investigation in this regard does suggest that an overall strengthening of the freestanding coating is indeed possible by increasing the strength of the intermediate layer by the above two methods. The detailed results of this study will be reported in a future publication. 5. Conclusions Room temperature tensile behaviour of a high activity Pt–aluminide coating has been evaluated in stand-alone condition using microtensile testing method. The coating had a typical threelayer microstructure. The tensile stress–strain plot was linear, indicating the coating to be brittle. Microstructural and fractographic evidence along with the fracture toughness values of the individual coating layers suggest that the failure of the microtensile sample occurred by initiation of a crack in the intermediate layer and its subsequent inside-out propagation. Acknowledgements The authors wish to acknowledge the assistance provided by the SMG, SFAG and EMG divisions of DMRL for carrying out the coating heat treatments and characterization studies. The authors are thankful to Director, DMRL for his permission to publish this research work. This research work has been funded by the Defence Research and Development Organization. References [1] R. Pichoir, in: D.R. Holmes, A. Rahmel (Eds.), Materials and Coatings to Resist High Temperature Corrosion, Applied Science Publishers, London, 1978, p. 271.
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[2] S. Bose, High Temperature Coatings, Butterworth-Heinemann Publishers, Oxford, 2007, p. 71. [3] G.W. Goward, D.H. Boone, Oxid. Metals 3 (1971) 475. [4] J. Angenete, K. Stiller, Surf. Coat. Technol. 150 (2002) 107. [5] A. Strang, E. Lang, in: R. Brunetand (Ed.), High Temperature Alloys for Gas Turbines, D. Reidel Publishing Company, Holland, 1982, p. 469. [6] K. Schneider, H.W. Grunling, Thin Solid Films 107 (1983) 395. [7] J.M. Veys, R. Mervel, Mater. Sci. Eng. 88 (1987) 253. [8] R. Nützel, E. Affeldt, M. Göken, Int. J. Fatigue 30 (2008) 313. [9] D. Pan, M.W. Chen, P.K. Wright, K.J. Hemker, Acta Mater. 51 (2003) 2205. [10] K.J. Hemker, W.N. Sharpe Jr., Annu. Rev. Mater. Res. 37 (2007) 93. [11] M. Eskner, R. Sandström, Surf. Coat. Technol. 165 (2003) 71.
[12] D.K. Das, V. Singh, S.V. Joshi, Metall. Mater. Trans. A 31A (2000) 2037. [13] D.K. Das, V. Singh, S.V. Joshi, Metall. Mater. Trans. A 29A (1998) 2173. [14] B. Gleeson, W. Wang, S. Hayashi, D.J. Sordelet, Mater. Sci. Forum 461–464 (2004) 213. [15] R. Villars, L.D. Calvert, Pearson’s Handbook of Crystallographic Data for Intermetallic Phases, vol. 4, second edn, ASM International, 1991, p. 4723. [16] D.B. Marshall, B.R. Lawn, A.G. Evans, J. Am. Ceram. Soc. 65 (11) (1982) 561. [17] R.D. Noebe, R.R. Bowman, M.V. Nathal, Int. Mater. Rev. 38 (4) (1993) 193. [18] D.B. Miracle, Acta Metall. Mater. 41 (3) (1993) 649. [19] C. Jiang, M.F. Besser, D.J. Sordelet, B. Gleeson, Acta Mater. 53 (2005) 2101. [20] D.K. Das, J.P. McDonald, S.M. Yalisove, T.M. Pollock, Spectrochem. Acta Part B 63 (2008) 27.
Contents lists available at ScienceDirect
Materials Science and Engineering A journal homepage: www.elsevier.com/locate/msea
Mechanism of failure in a free-standing Pt–aluminide bond coat during tensile testing at room temperature Md. Zafir Alam a,∗ , B. Srivathsa a , S.V. Kamat a , V. Jayaram b , N. Hazari a , D.K. Das a a b
Defence Metallurgical Research Laboratory, Hyderabad 500 058, India Indian Institute of Science, Bangalore 560 012, India
a r t i c l e
i n f o
Article history: Received 29 June 2009 Received in revised form 30 August 2009 Accepted 8 September 2009
Keywords: Bond coat Pt–aluminide Microtensile test Failure mechanism Crack nucleation and propagation Fracture toughness
a b s t r a c t The room temperature (RT) tensile behaviour of a free-standing high activity Pt–aluminide bond coat has been evaluated by microtensile testing technique. The coating had a typical three-layer microstructure. The stress–strain plot for the free-standing coating was linear, indicating the coating to be brittle at RT. Different fracture features were observed across the coating layers, namely quasi-cleavage in the outer layer and inner interdiffusion zone, and cleavage in the intermediate layer. By employing interrupted tensile test and observing the cross-sectional microstructure of the tested specimens, it was determined that failure of the microtensile samples occurred by the initiation of a single crack in the intermediate layer of the coating and its subsequent inside-out propagation. Such a mechanism of failure has been explained in terms of the fracture features observed across the sample thickness. This mechanism of failure is consistent with fracture toughness values of the individual coating layers. © 2009 Elsevier B.V. All rights reserved.
1. Introduction Nickel based superalloy components such as blades and vanes that operate in the hot sections of gas turbine engines are usually applied with Pt–aluminide (PtAl) bond coats for enhancing their high temperature oxidation resistance [1,2]. These coatings are formed on the superalloy first by depositing a layer of Pt and then subjecting the plated superalloy to a vacuum diffusion treatment. Subsequently, the superalloy is given an aluminizing treatment to form the coating. Depending on the aluminizing technique used, these coatings are classified as inward grown or outward grown [1,3,4]. These two varieties are also called as high activity and low activity coatings, respectively. The bond coats, which are typically 50–100 m thick, are based on the B2-NiAl intermetallic system. Therefore, they are inherently brittle and have a high brittle-to-ductile transition-temperature (BDTT) above 600 ◦ C [5–11]. PtAl coatings have been reported to degrade the mechanical properties, especially the tensile ductility, of the substrate superalloy at temperatures below the BDTT [5–7].
∗ Corresponding author at: Surface Engineering Group (Coatings), Defence Metallurgical Research Laboratory, P.O.-Kanchanbagh, Hyderabad 500 058, India. Tel.: +91 40 24586634; fax: +91 40 24340683. E-mail address: zafir [email protected] (Md. Z. Alam). 0921-5093/$ – see front matter © 2009 Elsevier B.V. All rights reserved. doi:10.1016/j.msea.2009.09.017
In most of the early literature, the effect of bond coats on the mechanical properties (including tensile) of the superalloys had been studied by conducting tests on the coated superalloy samples [5–8]. In recent years, however, it has been realized that it is important to understand the mechanical behaviour of the stand-alone coatings in order to gain better insight into their effect on the properties of the superalloys. Hemker and co-workers [9,10] have reported the tensile properties of outward grown (low activity) free-standing PtAl coating between room temperature and 1000 ◦ C. They have adopted microtensile testing technique to evaluate the tensile properties of these PtAl coatings. Their studies have established several important aspects of low activity PtAl coatings such as their BDTT being about 600 ◦ C. Above this temperature, ductility of the coating has been found to increase considerably while the strength drops very rapidly. Similar values of BDTT of PtAl coatings have also been evaluated by Eskner and Sandström [11] by using a different test method, i.e. miniaturized disk bending technique. Despite the above-mentioned reports on the tensile behaviour of free-standing PtAl coatings, to the best of our knowledge, there has not been any study correlating the microstructure of PtAl coatings with their tensile failure behaviour. In the present investigation, such a correlation has been made in case of a stand-alone high activity (inward grown) PtAl coating at RT. Based on this correlation, a mechanism of failure in the coating has been proposed. Microtensile test method has been adopted to determine the tensile properties of the coating.
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Fig. 1. Schematic of a microtensile sample. All dimensions are in mm.
2. Experimental details 2.1. Sample design Microtensile samples of the geometry as shown in Fig. 1, were used in the present study. The thickness of the sample, which was equal to the coating thickness was approximately 100 m. For the samples of the above geometry, conventional gripping of the specimen between two flat plattens invariably led to the breakage of the samples at the fillet region. This was caused due to the inherent brittleness of the coating. To overcome this problem, the samples were held in slotted grips for testing as shown in Fig. 2. In order to ensure the failure of the samples within the gage length, use of appropriate holding length, HL, and fillet radius, R, (Fig. 1) was found to be critical. To fix these two dimensions, a finite element method (FEM)-based elastic analysis was carried out. Using the appropriate boundary conditions imposed by the slotted grips, a combination of HL and R was determined such that, upon application of tensile loading, a uniform maximum stress was generated within the gage length while the elastic stress concentration at the fillet remained at a minimum value. It was determined that a combination of R = 0.5 mm and HL = 1 mm, which corresponded to an elastic stress concentration at the fillet of 1.02, would lead to failure of the microsamples within the gage length. Based on this analysis, samples with gage length (GL) of 2 mm, gage width (GW) of 0.5 mm, overall length of 8 mm, HL of 1 mm and R of 0.5 mm were prepared. Initial tensile tests indicated that the samples with the above dimensions failed within the gage length on most occasions. Therefore, these sample dimensions were used for evaluation of coating properties.
Fig. 2. A free-standing PtAl microtensile sample placed in the slotted grips while testing.
Fig. 3. Method adopted for the fabrication of microtensile specimens: (a) PtAl coated 0.5 mm thick superalloy strip, (b) wire EDM machining of samples from the coated strip and (c) a typical cut-out microtensile sample.
2.2. Sample preparation, tensile testing and microstructural characterization Directionally solidified (DS) rods of Ni-based superalloy CM247LC were used as the substrate material for developing the PtAl bond coats. The nominal composition (in wt.%) of the above alloy is 9.2 Co–8.1 Cr–9.5 W–5.6 Al–3.2 Ta–1.5 Hf–0.7 Ti–0.015 Zr–0.5 Mo–0.15 B–0.07 C–balance Ni. The coating was formed on 0.5 mm thick strips sliced from the above DS rods. For coating formation, a 5 m thick Pt layer was first electrodeposited on these strips. The plated strips were then subjected to a diffusion heat treatment in vacuum, following which they were given a pack aluminization treatment. A two-step high activity pack aluminizing process [1,2,12,13] was used in this study. Microtensile samples were cut from the coated strips by electro discharge machining (EDM) as shown in Fig. 3(b). The gage length of the microsamples was always parallel to the longitudinal axis of the strips, i.e. along 0 0 1 solidification direction (see Fig. 3(a)). The above cut-out samples (Fig. 3(c)) were then precisionpolished from one side using polishing papers of 600, 1000, 1500 and 2500 grades, respectively. Final finishing was carried out on polishing films containing 1 and 0.5 m diamond particles. The unpolished surface of the coating was comparatively much rougher having an average roughness (Ra ) of 1.3 m as compared to the 0.5 m finish of the polished surface (see Fig. 1). To prepare free-standing coating samples, polishing was carried out till the substrate was completely removed and the sample thickness became equal to the thickness of bond coat, which was approximately 100 m. A Walter–Bai Ag microtensile testing machine with 500 N loading capacity was employed for testing these freestanding coating samples at RT. Five coating microsamples were tested at an initial strain rate of 4.1 × 10−3 s−1 and a constant crosshead velocity of 0.5 mm/min. A pre-stress of 10 MPa was applied for proper gripping of the microsamples within the slotted grips (Fig. 2). Microstructural and fractography observations were made in a Quanta 400 scanning electron microscope (SEM) operating at 20 kV. A Cameca SX-100 electron probe micro analyzer (EPMA) operating at 20 kV was used for chemical analysis of the coating. Identification of the phases present in the coating was carried out by X-ray diffraction (XRD) method using a Philips PW-3020 X-ray diffractometer. Indentation studies were carried out using a Leica VHMT Auto microhardness tester and a CSM instrumented nanoindenter.
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Fig. 4. Microstructure of the as-formed coating showing the individual layers. Some of the voids present in the outer layer are indicated by arrows. Ni plating was given on the sample for edge-protection during metallographic polishing. 1 and 2 correspond to locations whose compositions have been indicated in the ternary phase diagram (Fig. 5).
3. Results
Fig. 6. Variation in Ni, Pt and Al concentrations across the coating thickness.
3.1. Coating microstructure The PtAl coating, as shown in Fig. 4, revealed a three-layer microstructure, typically obtained in a high activity aluminizing process [1,2,12,13]. The matrix phase in all the layers was B2NiAl having Pt in solid solution, i.e. B2-(Ni,Pt)Al. The outer layer additionally contained PtAl2 phase and numerous fine precipitates. The intermediate layer had a single-phase B2-(Ni,Pt)Al structure and contained comparatively much fewer precipitates. The B2(Ni,Pt)Al + PtAl2 structure of the outer layer and the B2-(Ni,Pt)Al structure of the intermediate layer (indicated by points 1 and 2, respectively, in Fig. 4) are consistent with the corresponding compositions of these layers, which can be confirmed based on the ternary phase diagram of Ni–Pt–Al reported by Gleeson et al. [14]. As shown in Fig. 5, the composition of a typical location in the outer layer (point 1 in Fig. 4) corresponds to the B2-(Ni,Pt)Al + PtAl2 phase field. Similarly, the composition of the intermediate layer (point 2 in Fig. 4) lies in the single-phase field of B2-(Ni,Pt)Al. The inner layer of the coating was the interdiffusion zone (IDZ), typically found in most diffusion aluminide coatings [1,2]. This layer is formed as a
result of the loss of Ni from the superalloy substrate required for coating formation [1,2,12,13]. Such loss of Ni leads to precipitation of numerous complex precipitates [1,2,12,13] as seen in the IDZ of the present coating (Fig. 4). The thicknesses of the abovementioned three layers were approximately 60, 25 and 20 m, respectively. The fine precipitates found in the outer layer of the coating were rich in Ni and W [13] and identified as orthorhombic NiW by XRD method [15]. A few irregular shaped Kirkendall voids were also observed in the top regions of the outer coating layer (Fig. 4). The coating did not contain any crack in as-coated condition. Being a diffusion coating, the present PtAl coating contained almost all the substrate alloying elements such as Cr, W, Ta, Hf, Mo and Co. The concentration profiles for Ni, Al and Pt, as shown in Fig. 6, revealed the graded composition of the coating. The Ni concentration increased continuously from about 25 at.% to nearly 50 at.% over the outer and the intermediate layers. Its concentration remained in the range 40–50 at.% in the IDZ. The Al concentration remained in the narrow range of 45–40 at.% over the outer and intermediate layers and fell to almost 20 at.% at the IDZ/substrate interface. The Pt concentration varied fairly sharply across the coating thickness with the outer layer, intermediate layer and the IDZ containing a maximum of about 10, 3 and 0.1 at.% Pt, respectively. 3.2. Tensile behaviour
Fig. 5. Isothermal section of Ni–Pt–Al ternary system at 1100 ◦ C [14]. Composition of locations 1 and 2 correspond to the compositions of regions marked 1 and 2 in Fig. 4.
The stress–strain plot for the free-standing PtAl coating, as shown in Fig. 7, was fairly linear and did not show any appreciable plastic yielding. The total fracture strain was determined to be about 0.2% and the plastic strain was negligibly small. Such low plasticity for the PtAl coating at RT is expected considering their high BDTT values, which are usually above 600 ◦ C [9,10]. The elastic modulus and the fracture stress were determined to be 100 ± 1 GPa and 215 ± 40 MPa, respectively. The linear stress–strain relation observed in the present study is similar to that reported for low activity PtAl coatings by Pan et al. [9]. The typical fracture surface of a failed coating microsample is presented in Fig. 8(a). As evident from this figure, the three coating layers showed distinctly different fracture features. The outer layer and the IDZ showed quasi-cleavage features (Fig. 8(b) and (d), respectively), while the intermediate layer exhibited cleavage features with river-bed patterns (see Fig. 8(c)).
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Fig. 7. Stress–strain plot for the free-standing PtAl coating.
3.3. Crack initiation and propagation After the completion of the tensile test, the failed pieces of the tensile sample were carefully examined for the presence of any crack on both as-coated and polished side (see Fig. 1) in SEM. Not even a single crack in the transverse direction, i.e. normal to the loading direction, could be found over the gage length on both the surfaces of the sample, as typically shown in Fig. 9(a) and (b). To confirm this fact further, metallographically polished longitudinal sections of the failed pieces (see Fig. 1) were also examined. No cracks (perpendicular to the coating thickness) in these sections could be located over the entire gage length, as evident from Fig. 9(c). Further, no crack could also be located on the edges on
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both polished and unpolished surfaces. This is an indication that cracking perpendicular to the loading direction did not occur on the thickness side of the microtensile sample. These observations clearly suggest that the failure of the sample during testing has most likely occurred by the generation and propagation of the first crack in the sample. To investigate this aspect further, interrupted tensile test on four different microtensile samples was conducted wherein the samples were loaded to stress levels of 50, 100, 150 and 180 MPa, respectively, and subsequently unloaded. Care was taken to prevent any breakage of the samples during removal from the slotted grips (Fig. 2). Subsequently, the longitudinal section of the samples was examined for the presence of any crack. The samples subjected to stress levels below 150 MPa did not exhibit any cracking through out the gage length. However, for samples subjected to stress levels of 150 MPa and beyond, only a single crack over the entire gage length was located as shown in Fig. 10(a) and (b). For the sample loaded to 150 MPa, the crack could be seen initiating in the intermediate layer of the coating (Fig. 10(a)). However, the sample loaded to a higher stress of 180 MPa revealed a crack that had propagated across the outer and the intermediate layers, and partially across IDZ (Fig. 10(b)). Even in this sample, the higher crack width in the intermediate layer of the coating suggests that the crack has initiated in this layer and propagated subsequently to the other two layers. The above observations can be summarized as follows. The failure of the stand-alone coating microtensile sample has occurred by the initiation and subsequent growth of the first crack in the sample. The crack has initiated in the precipitate-lean intermediate layer of the coating and subsequently propagated inside-out leading to the sample failure. In order to substantiate the above observation, the fracture toughness across the individual layers of the coating was estimated. For this estimation, a series of indentations were made in each of the coating layers using a Vickers indenter at various loads. Subsequently, they were observed in
Fig. 8. (a) Fracture surface observed on the failed microtensile sample showing the variation of fracture features across the coating thickness. Magnified view of the fracture surface in the outer layer (OL), intermediate layer (IL) and the IDZ are shown in (b), (c) and (d), respectively.
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Fig. 9. Micrographs depicting the various surfaces of the failed tensile sample: (a) as-coated top side, (b) polished side and, (c) longitudinal section of the gage. No crack was located on any of these regions of the sample.
Fig. 10. Longitudinal section of the coating microsamples subjected to interrupted tensile test in which the loading was stopped at (a) 150 MPa and (b) 180 MPa.
SEM to examine the presence of any cracking around the indentations. It was found that between 2 and 3 N loads, lateral cracks originated around the indentations. The minimum load at which cracking around the indentation was observed in any given layer was considered as the critical load for that layer. Since lateral cracking was observed around the indentations, the fracture toughness value was calculated from the relation proposed by Marshall et al. [16], as given in Eq. (1):
4 KIND =
Po A2o 2/3 (cot ) ıo
H4 E
(1)
In the above equation, KIND is the indentation fracture toughness, Po the threshold load for lateral cracking, Ao and ıo the dimensionless constants, the half included angle of the indenter, H the hardness of the material in GPa and E the indentation modulus in GPa. Using the appropriate values of the above-mentioned terms in the equation for each of the three coating layers, the fracture toughness for the individual layers was calculated, as presented in Table 1. It is evident that the fracture toughness of the intermediate layer at 1.15 MPa m1/2 is the lowest among the three coating layers, which supports the earlier observation of the crack initiating in the intermediate layer during tensile test.
4. Discussion As mentioned earlier, the general linear tensile behaviour of the present high activity PtAl coating (Fig. 7) was similar to that reported for low activity PtAl coating [9]. Further, the elastic modulus for the present coating at 100 ± 1 GPa was also similar to the value of 117 GPa reported by Pan et al. for the low activity coating [9]. However, the fracture strength of 215 MPa for the present coating was much higher than the value of 100 MPa reported for the low activity coatings [9]. The difference in the fracture strength values in the two studies can be ascribed to the difference in the stoichiometry and the associated defect strengthening
Table 1 Indentation fracture toughness, KIND of various layers of the Pt–aluminide coating. The data shown has been averaged over ten readings in each layer. The values of dimensionless constants Ao and ıo are 3/4 and 1200, respectively. The half included angle for the Vickers indenter was 68◦ . The values of H, E and KIND are the average of ten measurements.
Outer layer Intermediate layer IDZ
Po (N)
H (GPa)
E (GPa)
KIND (MPa m1/2 )
2 3 2
8.39 ± 0.64 4.81 ± 0.32 7.33 ± 0.17
302 ± 71 244 ± 28 411 ± 31
1.71 ± 0.13 1.15 ± 0.09 1.36 ± 0.02
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Fig. 11. Variation in the stoichiometry of the B2-(Ni,Pt)Al phase present in high activity and low activity coatings. The dashed line indicates stoichiometric (Ni,Pt)50 Al50 .
of the B2-NiAl phase present in the matrix in both the coatings. It is well known that the stoichiometric structure, i.e. Ni50 Al50 structure, is least defective in case of binary B2-NiAl phase [17,18]. On Al-rich side of the stoichiometry, i.e. for hyperstoichiometric structures, vacancy defects develop because of the unfilled Ni-sites. For hypostoichiometric (Ni-rich) structures, however, substitutional defects develop because of excess Ni atoms occupying Al lattice sites [17,18]. It has been widely reported that that the presence of such defects causes some degree of strengthening in bulk B2-NiAl. Further, vacancy defects in hyperstoichiometric B2-NiAl cause much greater strengthening as compared to the corresponding substitutional defects in hypostoichiometric B2 phase [17,18]. High activity coatings are generally much higher in Al content as compared to low activity coatings. Therefore, the B2 phase in the high activity coatings is expected to be largely hyperstoichiometric in structure. To examine this aspect in the present high activity coating, it was assumed that (Ni,Pt)50 Al50 is the stoichiometric composition for the B2 phase, since it is known that Pt occupies Ni-sites in the B2 lattice [19]. Assuming that the effect of site occupancy of other elements in the B2 phase can be neglected, the off-stoichiometric composition can be written as (Ni,Pt)50+d Al50−d , where d is a parameter (in at.%) representing the deviation from stoichiometry. The value of d is zero for stoichiometric composition, positive for hypostoichiometric composition and negative for hyperstoichiometric compositions. The value of d can be determined by equating (50 + d)/(50 − d) to the atom ratios of (Ni + Pt) to Al, as determined from the EPMA composition profiles (see Fig. 6). The variation of d corresponding to the single-phase B2-(Ni,Pt)Al across the thickness of the present coating has been plotted in Fig. 11. The B2-(Ni,Pt)Al + PtAl2 two phase region (approximately 25 m in thickness) of the outer layer has not been considered in the above plot (see Fig. 11) since it was impossible to accurately determine the composition of the B2 phase in this region using the EPMA. For comparison, the corresponding plot for a typical low activity coating (MDC150L) has also been given in the same figure. The plot for the MDC150L coating was carried out using the data reported by Das et al. [20]. It is clear from Fig. 11, that the B2 composition up to a depth of about 60 m in the present coating is hyperstoichiometric, whereas it is completely hypostoichiometric in case of MDC150L coating. The comparative higher fracture strength of 215 MPa for the present PtAl coating is most likely because of its hyperstoichiometric composition and the cor-
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responding high vacancy defects as compared to the low activity coating. However, since modulus is not affected significantly by the deviation from the stoichiometry [17,18], the modulus of the present coating was found to be similar. As previously mentioned, there was no crack present in the coating in the as-coated condition, although some voids in the outer layer could be observed (Fig. 4). Therefore, it can be concluded that the crack causing the failure of the microtensile sample has initiated during the tensile test. The microstructural evidence (Fig. 10(a) and (b)) as well as the fracture toughness data of the individual coating layers (Table 1) indicate that the above crack has initiated in the intermediate layer. The predominantly cleavage fracture in the intermediate layer of the coating (Fig. 8(c)) further supports the above conclusion. Thus, all the evidences, namely the origin of the crack as seen in the coating cross-section (Fig. 10(a)), fracture toughness values and the fracture surface features, together clearly suggest that the failure of the microtensile sample has occurred by initiation of a single crack in the intermediate layer which subsequently propagated inside-out (Fig. 10(b)). It is interesting that despite the presence of stress raisers such as voids close to the surface of the sample (Fig. 4) and also appreciably large roughness of 1.3 m on the unpolished side of the microtensile sample, the cracking initiated in the sub-surface intermediate layer (Fig. 10(a) and (b)). The first crack in the sample leading to failure clearly suggests that once a crack is initiated, its propagation is extremely easy in the brittle B2-(Ni,Pt)Al matrix of the coating. Thus, there is no apparent driving force for the generation of any additional crack in the sample during testing. From the above discussion, it appears that the overall strength of the coating would be primarily dependent on the strength of the intermediate layer. Therefore, strengthening of this layer would increase the overall strength of the coating. The strength of the intermediate layer can be enhanced through refining the grain size of the B2(Ni,Pt)Al phase constituting this layer. Increased Pt solid solution strengthening of the above phase may also be possible by enhancing the Pt content of the coating. In fact, our preliminary investigation in this regard does suggest that an overall strengthening of the freestanding coating is indeed possible by increasing the strength of the intermediate layer by the above two methods. The detailed results of this study will be reported in a future publication. 5. Conclusions Room temperature tensile behaviour of a high activity Pt–aluminide coating has been evaluated in stand-alone condition using microtensile testing method. The coating had a typical threelayer microstructure. The tensile stress–strain plot was linear, indicating the coating to be brittle. Microstructural and fractographic evidence along with the fracture toughness values of the individual coating layers suggest that the failure of the microtensile sample occurred by initiation of a crack in the intermediate layer and its subsequent inside-out propagation. Acknowledgements The authors wish to acknowledge the assistance provided by the SMG, SFAG and EMG divisions of DMRL for carrying out the coating heat treatments and characterization studies. The authors are thankful to Director, DMRL for his permission to publish this research work. This research work has been funded by the Defence Research and Development Organization. References [1] R. Pichoir, in: D.R. Holmes, A. Rahmel (Eds.), Materials and Coatings to Resist High Temperature Corrosion, Applied Science Publishers, London, 1978, p. 271.
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