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Microscale lamellar NiCoCrAlY coating with improved oxidation resistance
Surface & Coatings Technology 207 (2012) 110–116
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Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
Microscale lamellar NiCoCrAlY coating with improved oxidation resistance Hui Peng, Hongbo Guo ⁎, Jian He, Shengkai Gong ⁎ School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China
a r t i c l e
i n f o
Article history: Received 19 January 2012 Accepted in revised form 11 June 2012 Available online 16 June 2012 Keywords: Oxidation Alumina scale Coating Plasma activated electron beam-physical vapor deposition (PA EB-PVD)
a b s t r a c t A novel NiCoCrAlY coating was fabricated by plasma activated electron beam-physical vapor deposition (PA EB-PVD). Within the duplex coating, the β-NiAl phase precipitated as small lamellae (less than 1 μm thick) in a γ matrix, predominantly oriented perpendicular to the coating surface. The oxidation results at 1373 K show that the scaling rate constant during the steady-state oxidation of the novel coating is reduced by approximately 70% as compared with the conventional coating of similar composition produced by electron beam-physical vapor deposition (EB-PVD). This reduction in oxidation rate is attributed to a slower θ → α Al2O3 phase transformation during the initial stages of oxidation. © 2012 Elsevier B.V. All rights reserved.
1. Introduction MCrAlY (M = Ni, Co or Ni + Co) overlay coatings are commonly employed on superalloys or used as the bond coats (BCs) for zirconia based thermal barrier coating (TBC) systems [1–3]. The protection offered by such coatings against high-temperature oxidation and corrosion relies on the performance of the thermally grown oxide (TGO) developed between the ceramic topcoat and the BC [4–6]. Formation of a continuous and dense α-Al2O3 scale is preferred not only due to its high thermal and chemical stabilities, but also due to the low diffusivity of oxygen and metal ions in its hexagonal close-packed (HCP) structure [7]. Several attempts have been made to improve the high-temperature oxidation performance of the BCs [8,9], including varying their microstructure. The most common approach for surface modification is shot peening, which can cure the growth defects at the near surface region mechanically [10]. One further approach is to form a re-melted surface layer by pulsed electron beam (PEB) or laser treatment. This treatment of the BC was found to reduce the rate of alumina growth and to suppress the formation of Y-rich oxide pegs resulting in a smoother BC interface [11,12]. Reducing the coating's grain size is another effective way to improve the oxidation resistance. It was revealed that the selective formation of Al2O3 was a function of the coating grain size and that for the smallest grain size (~65 nm) a compact alumina scale was formed [13]. The aforementioned studies suggest that dense and refined microstructure, yielded either by controlling the deposition parameters or by post-treatment, is helpful to improve the oxidation resistance of MCrAlY coatings. ⁎ Corresponding authors. Tel.: + 86 10 8231 7117; fax: + 86 10 8233 8200. E-mail addresses: [email protected] (H. Guo), [email protected] (S. Gong). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.06.034
Recently, a double-layer NiCoCrAlY coating with improved oxidation resistance has been proposed [14]. The bottom layer produced by plasma activated electron beam-physical vapor deposition (PA EB-PVD) was featured with coarse-grained equiaxed microstructure, which was different from the columnar microstructure of the top layer produced by electron beam-physical vapor deposition (EB-PVD). The coarsegrained microstructure reduced the density of grain boundaries, which served as high-diffusivity paths for the outward diffusion of refractory elements in the substrate. Thus it resulted in an improved oxidation resistance in the double-layer coating as compared with the conventional single-layer coating. Actually, the plasma assistance was first used to produce leader free MCrAlY coatings, aiming to eliminate the linear defects [15]. After that, coatings of lamellar phase microstructure were also fabricated by using this technique, but only briefly documented [16,17]. In this work, a lamellar structured NiCoCrAlY coating was fabricated. A detailed study of the coating microstructure was performed. The improved oxidation resistance of the coating and possible oxidation mechanisms were discussed. 2. Experimental 2.1. Materials and sample preparation Disk-shaped specimens cut from a cast alloy ingot were used as substrates. The alloy ingot was produced by vacuum induction melting, with the nominal composition of Ni–20Co–22Cr–8.8Al–1.5Y (in wt.%). The use of NiCoCrAlY substrates allowed accessing the oxidation behavior of coatings independently, minimizing the influence of element inter-diffusion. For instance, it has been widely reported that the outward diffusion of refractory elements from the superalloy can lead to premature failure of the thermally grown oxide (TGO) as they diffuse into the coating [18]. Prior to coating deposition, the substrate surfaces
H. Peng et al. / Surface & Coatings Technology 207 (2012) 110–116
were ground with 800 grit SiC paper, ultrasonically cleaned in acetone, and thoroughly dried. Two kinds of NiCoCrAlY coatings were deposited onto the prepared substrates: (1) a conventional NiCoCrAlY coating, denoted as EB coating, for the comparative study; (2) a novel structured NiCoCrAlY coating, denoted as PA coating. The EB coating was produced by evaporation of the above mentioned alloy ingot with the conventional EB-PVD process. By controlling the deposition parameters, the EB coating was obtained with the similar composition to that of the evaporating ingot. On the other hand, the PA coating was produced by PA EB-PVD, the principle and deposition parameters of which were already shown elsewhere [14]. It is worth noting that the ingot for producing the PA coating was predesigned with a different composition of Ni–20Co–18Cr–12.3Al–1.5Y (wt.%), containing less Cr and more Al, in order to obtain the desired coating composition similar to the EB coating. This compositional deviation has been frequently observed in compound coatings grown by plasma assisted deposition [19,20]. After deposition, all the coating specimens were annealed in vacuum at 1323 K for 4 h, with a vacuum pressure of ~3×10− 3 Pa, in order to achieve homogenization of the coatings.
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Prior to and after oxidation, a variety of analysis techniques was employed to study the coating and oxide composition and microstructure. The microstructure of the coatings was characterized by a scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS) and back scattering electron (BSE) detector. The average composition and the composition of the γ and β phases in the coatings were determined by an electron probe micro-analyzer (EPMA). The volume fraction of β phase within the coatings was quantified by an image analysis software (Adobe® Photoshop® CS3). The phases of TGO formed on the coating surface after the TGA experiments were identified by Xray diffraction (XRD) using Cu Kα radiation. The crystallographic forms of the alumina scale after short-term oxidation tests were characterized by photostimulated-luminescence spectroscopy. After short-term oxidation tests for 5 h and TGA oxidation tests, coating specimens were fractured in liquid nitrogen for cross-sectional investigation of the oxide scale. 3. Results and discussion 3.1. Microstructure of the PA EB-PVD coating
2.2. Oxidation tests and analysis Static oxidation tests were performed using a thermogravimetric analyzer (TGA, Cahn Thermax 700) at a temperature of 1373 K for 100 h. An air flow of 80 ml/min was passed through the TGA at a pressure of 10 5 Pa. In order to ascertain the reproducibility, three experimental runs for each coating were performed. Weight measurements were performed every 1 s. Moreover, in order to study the oxidation mechanisms in more detail, additional short-term static oxidation tests were performed in a horizontal alumina furnace. For each coating, specimens were oxidized for 10 min, 1 h, 5 h and 10 h at 1373 K in air. Prior to introduction of the specimens, the furnace tube was heated to a temperature of 1373 K. Next, the specimens were placed in annealed alumina crucibles and transported pneumatically to the hot zone of the furnace within several seconds. After oxidation, the specimens were transported quickly out of the furnace and cooled to room temperature.
Fig. 1a and b compares the SEM surface morphologies of the asdeposited coatings. The EB coating exhibited a typical columnar microstructure with grain size of ~3 μm. Columnar gaps (indicated by white arrows in Fig. 1a) and a large amount of fine pores were also observed. For the PA coating, as shown in Fig. 1b, the grain boundaries became less clear. The coating surface was much denser and smoother than that of the EB coating. The surface of each grain in the PA coating had a terraced substructure, which was consistent with the growth terraces observed in a NiAl coating prepared by electron beam-directed vapor deposition (EB-DVD) [21]. The formation of the terraced surface was caused by the enhanced adatom mobility induced by plasma activation and substrate bias. Similar to the NiCoCrAlY coating reported in Ref. [14], it can also be inferred that a transition from columnar to equiaxed grains had occurred in the PA coating. Both the coatings had a similar average composition (Table 1) and a duplex phase microstructure consisting of precipitates of the Al-rich
Fig. 1. SEM images of the as-deposited coatings: surface morphologies for the EB coating (a) and the PA coating (b), and BSE cross-sections for the EB coating (c) and the PA coating (d).
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Table 1 Chemical compositions of the EB and PA coatings measured by EPMA (in wt.%) as well as the volume percentage of β precipitates identified by Adobe® Photoshop® CS3. Composition
EB NiCoCrAlY
PA NiCoCrAlY
Average composition Composition γ Composition β Volume β
49.2Ni–22.5Co–19.4Cr–8.9Al– 1.5Y 45.8Ni–24.9Co–21.2Cr–8.0Al 59.5Ni–15.1Co–7.9Cr–15.5Al 17%
46.0Ni–24.5Co–20.4Cr–9.1Al– 1.5Y 40.1Ni–31.3Co–24.5Cr–4.1Al 57.7Ni–14.6Co–8.9Cr–18.8Al 61%
β phase in the matrix of the Al-poor γ phase. The EB and PA coatings differed mainly with respect to the quantity and distribution of β phase within the coating. The β phase precipitated as islands within the EB coating (Fig. 1c), whereas in the PA coating, the β phase was shaped as well-arranged lamellae, with thickness of less than 1 μm, predominantly oriented perpendicular to the coating surface (Fig. 1d). For both coatings, as shown in Table 1, (1) the PA coating had a larger volume percentage and smaller size of β phase, (2) the PA coating contained less Al and more Co and Cr in its γ phase. It could also be clearly observed in Fig. 1c that the Y in the EB coating predominantly precipitated as NiY-rich phases, situated along the γ/β phase and γ/γ′ grain boundaries [12,22]. In the case of the PA coating, although the subsequent XRD detection revealed the existence of Y-rich pegs, it was hard detecting where the Y was situated. This might be caused by the formation of plenty of γ/β phase boundaries, along which the Y was distributed more homogeneously. The significant differences in the distribution and composition of the phases between the two coatings were the result of the different deposition conditions of EB-PVD and PA EB-PVD processes. For the PA coating, the formation of the finely spaced lamellar microstructure was attributed to two possible reasons: (1) a strongly textured γ matrix was obtained in the PA coating, see Fig. 3 and Ref. [14]. The preferred orientation of the γ matrix promoted directional growth of the β precipitants. (2) The higher adatom mobility provided the driving force for the formation of plenty of γ/β phase boundaries. The allocation of elements between the γ and β phases was in line with the increase of the β phase volume fraction. Further investigations are needed to understand the growth mechanism of the PA coating.
Fig. 3. XRD patterns of the EB and PA coatings after static oxidation at 1373 K for 100 h.
~ 1.3 mg cm − 2 after 100 h of oxidation, while the mass gain for the PA coating is ~ 0.8 mg cm − 2 after the same oxidation period (Fig. 2a). This indicated that the PA coating exhibited a lower oxidation rate than the EB coating. The calculated oxidation parabolic rate constants for the EB coating and PA coating were 2.9×10−6 mg2 cm−4 s−1 and 8.0×10−7 mg2 cm−4 s−1, respectively (Fig. 2b). Compared to the oxidation of the EB coating, the scaling rate constant of the PA coating was reduced by approximately 70%. Fig. 3 shows the XRD patterns of the coating specimens after the 100 h TGA experiments. Similar oxides were detected on both coatings. α-Al2O3 was formed as the main oxide. No spinel phase but YAlOx phase was present. The surface morphologies of the oxidized coatings are shown in Fig. 4. It can clearly be seen that the alumina scales formed on the EB and PA coatings exhibited distinctly different morphologies. For the EB coating, the scale was grown with a lace-like network morphology (Fig. 4a), implying that a preferential diffusion of oxygen from oxide surface toward oxide/coating interface took place during the
3.2. TGA experiments The representative TGA curves show that the oxidation of both coatings at 1373 K followed two oxidation stages: a fast initial oxidation stage, followed by the second stage of steady-state parabolic growth (Fig. 2). The EB coating yielded a mass gain of
Fig. 2. Oxidation kinetics of the EB and PA coatings in air at 1373 K for 100 h.
Fig. 4. SEM images for the surfaces of the EB and PA coatings after oxidation at 1373 K for 100 h: (a) EB coating; (b) PA coating.
H. Peng et al. / Surface & Coatings Technology 207 (2012) 110–116
high-temperature oxidation [23]. On the contrary, the oxidized PA coating was covered with numerous blunt needlelike oxides (Fig. 4b).
3.3. Short-term oxidation tests It has been well recognized that the TGO formed at the initial oxidation stage plays an important role on its subsequent thickening [24]. Therefore, short-term oxidation tests for different periods were performed to investigate the scale growth at the initial oxidation stage. Four time points were chosen based on the slope change of the TGA plot: after 10 min of oxidation, i.e. during the rapid increase in mass gain; after 1 h and 5 h of oxidation, i.e. during the time period in which the mass gain dropped drastically; and after 10 h, i.e. during the steady-state period of oxidation.
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Fig. 5 shows the surface morphologies of the coating specimens after different periods of oxidation. From Fig. 5a and b, it can be observed that after oxidation for 10 min at 1373 K, oxide was formed on the surface of both coatings. Lace-like oxides and whisker-like oxides were developed on the surface of the EB coating and PA coating, respectively. As the oxidation continued, the scale morphology didn't change significantly for the EB coating (Fig. 5c, e and g). On the other hand, the scale formed on the PA coating after different oxidation times exhibited an obvious morphology evolution. After oxidation for 10 min, whisker-like oxides partially covered the coating surface (Fig. 5b). With continued oxidation, the scale maintained the whisker-like morphology but grew denser and denser until 5 h (Fig. 5d and f). By 10 h oxidation, the whisker-like crystals became blunt, although the scale kept on thickening (Fig. 5h). Examination of the fractured sections revealed that the scale for the EB coating after 5 h of oxidation was composed largely of equiaxed grains
Fig. 5. SEM images for the surfaces of the EB and PA coatings after short-term oxidation tests for 10 min, 1 h, 5 h and 10 h at 1373 K.
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(Fig. 6a), whereas on the PA coating it mainly consisted of θ-Al2O3 whiskers (Fig. 6b). After 100 h of oxidation, although growing thicker (with a thickness of ~6.5 μm), the scale on the EB coating still exhibited an equiaxed structure. On the other hand, the θ-Al2O3 whiskers on the PA coating became blunt with continued oxidation, showing a twolayered structure: an ~1 μm thick outer layer consisting of equiaxed grains and a 1.2 μm thick inner layer comprising of columnar grains (Fig. 6d). The phase evolution of alumina scales formed on the EB and PA coating surface with oxidation time at 1373 K is shown in Fig. 7, where the luminescence intensity is plotted against the wavenumber. As shown in Fig. 7a, θ-Al2O3 and α-Al2O3 signals were detected in the EB coating after 10 min of oxidation. With continued oxidation, the luminescence from the θ-Al2O3 decreased in intensity and by 5 h was barely discernible, indicating almost complete transformation to the stable α-Al2O3 phase. In the case of PA coating, only θ-Al2O3 but no α-Al2O3 was detected in the very early stage (10 min). Although the subsequent phase transformation sequence was similar to that of the EB coating, the times for the transformations were significantly longer. It took ~ 10 h to accomplish the θ → α transformation compared with ~5 h for the EB coating. 4. Discussion 4.1. Effect of coating microstructure on oxidation resistance In most cases, the microstructure of the oxide scale is a decisive factor for the oxidation resistance of the coating [25–27]. During the initial oxidation stage, the fast growing θ-Al2O3 that formed on both coatings led to a rapid mass gain. However, the θ → α transformation, which was detected after 1 h of oxidation, lowered down the scaling rate. It has been reported that the initial θ → α phase transformation occurs at the oxide/metal interface [28]. Once a continuous α-Al2O3 layer is formed, the subsequent oxidation will be controlled by diffusion through the α-Al2O3 layer. This was supported by the observation of the drastic drop of the scaling rate after ~1 h of oxidation for both
coatings. For the EB coating, a faster θ → α transformation occurred, but this fast transformation was accompanied with the formation of cracks in the oxide scale, as shown in Fig. 6a (indicated by white arrow). These cracks were the result of the development of tensile stress in the oxide scale, induced by the ~14% volume reduction associated with the θ → α transformation. The presence of micro-cracks allowed for the fast growth of new oxides and consequently accelerated the oxidation rate [26]. In contrast, a slower θ→α transition could lead to an extended lifetime of the θ phase layer, which was proposed to be beneficial to the relaxation of the stresses created by the growth of α-Al2O3. During the steady-state oxidation stage, the PA coating exhibited a better oxidation resistance compared with the EB coating. Former researchers have reported that a slow-growing and adherent α-Al2O3 layer on the coating surface can be promoted by performing a preoxidation treatment under a low-oxygen partial pressure [24,27]. The explanation in the work for the better oxidation resistance after pre-oxidation was that if θ-Al2O3 was well developed in the scale, the finally obtained α-Al2O3 grains would be large and thus decrease the oxidation rate, since the scale grew predominantly by the diffusion of oxygen along the α-Al2O3 grain boundaries. A similar explanation applied to the oxidation behavior of the PA coating in the present work, except that, for the PA coating studied here, the formation of metastable θ-Al2O3 was promoted even at atmospheric air pressure. Accordingly, the lower oxidation rate for the PA coating was attributed to the formation of a more protective scale with larger α-Al2O3 grains, as compared with the scale composed of quickly transformed, fine α-Al2O3 grains. It was also noted from the XRD results (Fig. 3) that a higher amount of YAlOx was formed in the EB coating after 100 h of oxidation. Since oxygen diffuses faster through YAlOx than through alumina, the increased amount of YAlOx in the scale also contributed to the higher oxidation rate of the EB coating. This can be interpreted that the columnar grain boundaries in the EB coating provided plenty of high-diffusivity paths for the outward diffusion of Y. In such a case, Y migrated rapidly along the grain boundaries to the coating surface and thickened the scale. Within the PA coating, the equiaxed grain
Fig. 6. SEM images for the fractured sections of the EB and PA coatings after short-term oxidation tests for 5 h and TGA experiments for 100 h at 1373 K. (Note that the images are shown at different magnifications.)
H. Peng et al. / Surface & Coatings Technology 207 (2012) 110–116
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Fig. 7. Evolution of luminescence spectrum for the EB and PA coatings after short-term oxidation tests at 1373 K: (a) EB coating; (b) PA coating.
reduced the outward diffusion rate of Y in some degree, then less YAlOx was formed. Similar results and discussions have been reported in Ref. [14].
4.2. Effect of coating microstructure on TGO formation The explanation given above raises the following question: why was the formation of θ-Al2O3 greatly enhanced in the PA coating? It is known that certain alloying elements can affect the kinetics of the θ→α transformation (for instance, Cr accelerates the θ→α transformation and Y retards the θ→α transformation) [25]. However, it can be noted that the EB and PA coatings had similar average compositions, thus the element competition effect can be eliminated. Therefore, it can be inferred that it was the coating microstructure that played the decisive role in the oxidation behavior. Another retarded θ → α transformation was observed at local regions of a plasma sprayed NiCoCrAlY coating [29]. A conclusion was drawn that α-Al2O3 and θ-Al2O3 initially formed on different phases in the coating and that there was then a slow transformation during oxidation of the θ-Al2O3 phase to α-Al2O3 correlated to the β-NiAl phase. Contrary to this, a homogeneous θ-Al2O3 scale was developed initially on the PA coating. The micro-lamella β phase distributed in the γ matrix within the PA coating was more beneficial in promoting the outward diffusion of Al to the coating surface. This is attributed to the following two reasons: (1) the larger volume percentage and smaller size of the β phase result in a higher density of γ/β phase boundaries in the coating, thereby providing fast diffusion paths for Al toward the coating surface [12,30], and (2) the nucleation of θ-Al2O3 is favored on the β phase of the coating surface due to preferred (epitaxial) orientation relationships between θ-Al2O3 and β-NiAl [31]. The finely spaced γ/β phase promoted the homogeneous growth of θ-Al2O3.
5. Conclusions A novel structured NiCoCrAlY coating was produced by plasma activated electron beam-physical vapor deposition (PA EB-PVD). Within the coating, the well-arranged β-NiAl phase precipitated as small lamellae (less than 1 μm in thickness) in the γ matrix, predominantly
oriented perpendicular to the coating surface. The conclusions drawn from the oxidation investigations are given as follows: (1) The scaling rate constant (8.0 × 10 − 7 mg 2 cm − 4 s − 1) during the steady-state of the novel coating was reduced by approximately 70% as compared with the conventional coating (2.9 × 10 − 6 mg 2 cm − 4 s − 1) of similar composition produced by electron beam-physical vapor deposition (EB-PVD). (2) The vertical arrangement of the micro-lamellae in the PA coating that provided fast diffusion paths promoted the outward diffusion of Al cations and the subsequent formation of well developed θ-Al2O3 scales. (3) The α-Al2O3 that transformed from the θ-Al2O3 grains grew at a much slower rate during the steady-state oxidation than the scale developed on a conventional NiCoCrAlY coating deposited by EB-PVD. Acknowledgments This work was financially supported by the National Basic Research Program (973 Program) of China under grant No. 2010CB631200 and the National Nature Science Foundations of China (NSFC, No.50771009, No. 50731001 and No. 51071013). References [1] P. Richer, M. Yandouzi, L. Beauvais, B. Jodoin, Surf. Coat. Technol. 204 (2010) 3962. [2] Y. Li, C.J. Li, G.J. Yang, L.K. Xing, Surf. Coat. Technol. 205 (2010) 2225. [3] S.M. Jiang, C.Z. Xu, H.Q. Li, J. Ma, J. Gong, C. Sun, Corros. Sci. 52 (2011) 1746. [4] D. Naumenko, V. Shemet, L. Singheiser, W.J. Quadakkers, J. Mater. Sci. 44 (2009) 1687. [5] J. Ma, S.M. Jiang, H.Q. Li, W.X. Wang, J. Gong, C. Sun, Corros. Sci. 53 (2011) 1417. [6] L.C. Chen, C. Zhang, Z.G. Yang, Corros. Sci. 53 (2011) 374. [7] B. Gleeson, N. Mu, S. Hayashi, J. Mater. Sci. 44 (2009) 1704. [8] H. Guo, Y. Cui, H. Peng, S. Gong, Corros. Sci. 52 (2011) 1440. [9] H. Lan, Z.G. Yang, Z.X. Xia, Y.D. Zhang, C. Zhang, Corros. Sci. 53 (2011) 1476. [10] J.F. Loersch, J.W. Neal, U. S. Pat. 4514469, (1981). [11] R.G. Wellman, A. Scrivani, G. Rizzi, A. Weisenburger, F.H. Tenailleau, J.R. Nicholls, Surf. Coat. Technol. 202 (2007) 709. [12] T.J. Nijdam, C. Kwakernaak, W.G. Sloof, Metall. Mater. Trans. A 37 (2006) 683. [13] Z. Liu, W. Gao, K. Dahm, F. Wang, Scr. Mater. 37 (1997) 1551. [14] H. Peng, H. Guo, J. He, S. Gong, Surf. Coat. Technol. 205 (2011) 4658. [15] H.A. Beale, T.E. Strangman, E.W. Taylor, U. S. Pat. 4109061, (1978).
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[16] Y.T. Lee, K. Fritscher, In: in: G.E. Broszeit, H. Exner (Eds.), Advanced Materials and Processes—Proceedings of the First European Conference, Aachen, Germany, 1989, p. 1255. [17] U. Schulz, K. Fritscher, In: in: K. Kokini (Ed.), Ceramic Coatings: ASME Winter Annual Meeting, New Orleans, Louisiana, 1993, p. 163. [18] J. Muller, D. Neuschutz, Vacuum 71 (2003) 247. [19] S. PalDey, S.C. Deevi, Mater. Sci. Eng., A 361 (2003) 1. [20] P.A. Lindfors, W.M. Mularie, G.K. Wehner, Surf. Coat. Technol. 29 (1986) 275. [21] Z. Yu, D.D. Hass, H.N.G. Wadley, Mater. Sci. Eng., A 394 (2005) 43. [22] T.J. Nijdam, L.P.H. Jeurgens, J.H. Chen, W.G. Sloof, Oxid. Met. 64 (2005) 355. [23] R. Prescott, M.J. Graham, Oxid. Met. 38 (1992) 233.
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Contents lists available at SciVerse ScienceDirect
Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
Microscale lamellar NiCoCrAlY coating with improved oxidation resistance Hui Peng, Hongbo Guo ⁎, Jian He, Shengkai Gong ⁎ School of Materials Science and Engineering, Beihang University (BUAA), No. 37 Xueyuan Road, Beijing, 100191, China
a r t i c l e
i n f o
Article history: Received 19 January 2012 Accepted in revised form 11 June 2012 Available online 16 June 2012 Keywords: Oxidation Alumina scale Coating Plasma activated electron beam-physical vapor deposition (PA EB-PVD)
a b s t r a c t A novel NiCoCrAlY coating was fabricated by plasma activated electron beam-physical vapor deposition (PA EB-PVD). Within the duplex coating, the β-NiAl phase precipitated as small lamellae (less than 1 μm thick) in a γ matrix, predominantly oriented perpendicular to the coating surface. The oxidation results at 1373 K show that the scaling rate constant during the steady-state oxidation of the novel coating is reduced by approximately 70% as compared with the conventional coating of similar composition produced by electron beam-physical vapor deposition (EB-PVD). This reduction in oxidation rate is attributed to a slower θ → α Al2O3 phase transformation during the initial stages of oxidation. © 2012 Elsevier B.V. All rights reserved.
1. Introduction MCrAlY (M = Ni, Co or Ni + Co) overlay coatings are commonly employed on superalloys or used as the bond coats (BCs) for zirconia based thermal barrier coating (TBC) systems [1–3]. The protection offered by such coatings against high-temperature oxidation and corrosion relies on the performance of the thermally grown oxide (TGO) developed between the ceramic topcoat and the BC [4–6]. Formation of a continuous and dense α-Al2O3 scale is preferred not only due to its high thermal and chemical stabilities, but also due to the low diffusivity of oxygen and metal ions in its hexagonal close-packed (HCP) structure [7]. Several attempts have been made to improve the high-temperature oxidation performance of the BCs [8,9], including varying their microstructure. The most common approach for surface modification is shot peening, which can cure the growth defects at the near surface region mechanically [10]. One further approach is to form a re-melted surface layer by pulsed electron beam (PEB) or laser treatment. This treatment of the BC was found to reduce the rate of alumina growth and to suppress the formation of Y-rich oxide pegs resulting in a smoother BC interface [11,12]. Reducing the coating's grain size is another effective way to improve the oxidation resistance. It was revealed that the selective formation of Al2O3 was a function of the coating grain size and that for the smallest grain size (~65 nm) a compact alumina scale was formed [13]. The aforementioned studies suggest that dense and refined microstructure, yielded either by controlling the deposition parameters or by post-treatment, is helpful to improve the oxidation resistance of MCrAlY coatings. ⁎ Corresponding authors. Tel.: + 86 10 8231 7117; fax: + 86 10 8233 8200. E-mail addresses: [email protected] (H. Guo), [email protected] (S. Gong). 0257-8972/$ – see front matter © 2012 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2012.06.034
Recently, a double-layer NiCoCrAlY coating with improved oxidation resistance has been proposed [14]. The bottom layer produced by plasma activated electron beam-physical vapor deposition (PA EB-PVD) was featured with coarse-grained equiaxed microstructure, which was different from the columnar microstructure of the top layer produced by electron beam-physical vapor deposition (EB-PVD). The coarsegrained microstructure reduced the density of grain boundaries, which served as high-diffusivity paths for the outward diffusion of refractory elements in the substrate. Thus it resulted in an improved oxidation resistance in the double-layer coating as compared with the conventional single-layer coating. Actually, the plasma assistance was first used to produce leader free MCrAlY coatings, aiming to eliminate the linear defects [15]. After that, coatings of lamellar phase microstructure were also fabricated by using this technique, but only briefly documented [16,17]. In this work, a lamellar structured NiCoCrAlY coating was fabricated. A detailed study of the coating microstructure was performed. The improved oxidation resistance of the coating and possible oxidation mechanisms were discussed. 2. Experimental 2.1. Materials and sample preparation Disk-shaped specimens cut from a cast alloy ingot were used as substrates. The alloy ingot was produced by vacuum induction melting, with the nominal composition of Ni–20Co–22Cr–8.8Al–1.5Y (in wt.%). The use of NiCoCrAlY substrates allowed accessing the oxidation behavior of coatings independently, minimizing the influence of element inter-diffusion. For instance, it has been widely reported that the outward diffusion of refractory elements from the superalloy can lead to premature failure of the thermally grown oxide (TGO) as they diffuse into the coating [18]. Prior to coating deposition, the substrate surfaces
H. Peng et al. / Surface & Coatings Technology 207 (2012) 110–116
were ground with 800 grit SiC paper, ultrasonically cleaned in acetone, and thoroughly dried. Two kinds of NiCoCrAlY coatings were deposited onto the prepared substrates: (1) a conventional NiCoCrAlY coating, denoted as EB coating, for the comparative study; (2) a novel structured NiCoCrAlY coating, denoted as PA coating. The EB coating was produced by evaporation of the above mentioned alloy ingot with the conventional EB-PVD process. By controlling the deposition parameters, the EB coating was obtained with the similar composition to that of the evaporating ingot. On the other hand, the PA coating was produced by PA EB-PVD, the principle and deposition parameters of which were already shown elsewhere [14]. It is worth noting that the ingot for producing the PA coating was predesigned with a different composition of Ni–20Co–18Cr–12.3Al–1.5Y (wt.%), containing less Cr and more Al, in order to obtain the desired coating composition similar to the EB coating. This compositional deviation has been frequently observed in compound coatings grown by plasma assisted deposition [19,20]. After deposition, all the coating specimens were annealed in vacuum at 1323 K for 4 h, with a vacuum pressure of ~3×10− 3 Pa, in order to achieve homogenization of the coatings.
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Prior to and after oxidation, a variety of analysis techniques was employed to study the coating and oxide composition and microstructure. The microstructure of the coatings was characterized by a scanning electron microscope (SEM) equipped with energy dispersive spectroscopy (EDS) and back scattering electron (BSE) detector. The average composition and the composition of the γ and β phases in the coatings were determined by an electron probe micro-analyzer (EPMA). The volume fraction of β phase within the coatings was quantified by an image analysis software (Adobe® Photoshop® CS3). The phases of TGO formed on the coating surface after the TGA experiments were identified by Xray diffraction (XRD) using Cu Kα radiation. The crystallographic forms of the alumina scale after short-term oxidation tests were characterized by photostimulated-luminescence spectroscopy. After short-term oxidation tests for 5 h and TGA oxidation tests, coating specimens were fractured in liquid nitrogen for cross-sectional investigation of the oxide scale. 3. Results and discussion 3.1. Microstructure of the PA EB-PVD coating
2.2. Oxidation tests and analysis Static oxidation tests were performed using a thermogravimetric analyzer (TGA, Cahn Thermax 700) at a temperature of 1373 K for 100 h. An air flow of 80 ml/min was passed through the TGA at a pressure of 10 5 Pa. In order to ascertain the reproducibility, three experimental runs for each coating were performed. Weight measurements were performed every 1 s. Moreover, in order to study the oxidation mechanisms in more detail, additional short-term static oxidation tests were performed in a horizontal alumina furnace. For each coating, specimens were oxidized for 10 min, 1 h, 5 h and 10 h at 1373 K in air. Prior to introduction of the specimens, the furnace tube was heated to a temperature of 1373 K. Next, the specimens were placed in annealed alumina crucibles and transported pneumatically to the hot zone of the furnace within several seconds. After oxidation, the specimens were transported quickly out of the furnace and cooled to room temperature.
Fig. 1a and b compares the SEM surface morphologies of the asdeposited coatings. The EB coating exhibited a typical columnar microstructure with grain size of ~3 μm. Columnar gaps (indicated by white arrows in Fig. 1a) and a large amount of fine pores were also observed. For the PA coating, as shown in Fig. 1b, the grain boundaries became less clear. The coating surface was much denser and smoother than that of the EB coating. The surface of each grain in the PA coating had a terraced substructure, which was consistent with the growth terraces observed in a NiAl coating prepared by electron beam-directed vapor deposition (EB-DVD) [21]. The formation of the terraced surface was caused by the enhanced adatom mobility induced by plasma activation and substrate bias. Similar to the NiCoCrAlY coating reported in Ref. [14], it can also be inferred that a transition from columnar to equiaxed grains had occurred in the PA coating. Both the coatings had a similar average composition (Table 1) and a duplex phase microstructure consisting of precipitates of the Al-rich
Fig. 1. SEM images of the as-deposited coatings: surface morphologies for the EB coating (a) and the PA coating (b), and BSE cross-sections for the EB coating (c) and the PA coating (d).
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Table 1 Chemical compositions of the EB and PA coatings measured by EPMA (in wt.%) as well as the volume percentage of β precipitates identified by Adobe® Photoshop® CS3. Composition
EB NiCoCrAlY
PA NiCoCrAlY
Average composition Composition γ Composition β Volume β
49.2Ni–22.5Co–19.4Cr–8.9Al– 1.5Y 45.8Ni–24.9Co–21.2Cr–8.0Al 59.5Ni–15.1Co–7.9Cr–15.5Al 17%
46.0Ni–24.5Co–20.4Cr–9.1Al– 1.5Y 40.1Ni–31.3Co–24.5Cr–4.1Al 57.7Ni–14.6Co–8.9Cr–18.8Al 61%
β phase in the matrix of the Al-poor γ phase. The EB and PA coatings differed mainly with respect to the quantity and distribution of β phase within the coating. The β phase precipitated as islands within the EB coating (Fig. 1c), whereas in the PA coating, the β phase was shaped as well-arranged lamellae, with thickness of less than 1 μm, predominantly oriented perpendicular to the coating surface (Fig. 1d). For both coatings, as shown in Table 1, (1) the PA coating had a larger volume percentage and smaller size of β phase, (2) the PA coating contained less Al and more Co and Cr in its γ phase. It could also be clearly observed in Fig. 1c that the Y in the EB coating predominantly precipitated as NiY-rich phases, situated along the γ/β phase and γ/γ′ grain boundaries [12,22]. In the case of the PA coating, although the subsequent XRD detection revealed the existence of Y-rich pegs, it was hard detecting where the Y was situated. This might be caused by the formation of plenty of γ/β phase boundaries, along which the Y was distributed more homogeneously. The significant differences in the distribution and composition of the phases between the two coatings were the result of the different deposition conditions of EB-PVD and PA EB-PVD processes. For the PA coating, the formation of the finely spaced lamellar microstructure was attributed to two possible reasons: (1) a strongly textured γ matrix was obtained in the PA coating, see Fig. 3 and Ref. [14]. The preferred orientation of the γ matrix promoted directional growth of the β precipitants. (2) The higher adatom mobility provided the driving force for the formation of plenty of γ/β phase boundaries. The allocation of elements between the γ and β phases was in line with the increase of the β phase volume fraction. Further investigations are needed to understand the growth mechanism of the PA coating.
Fig. 3. XRD patterns of the EB and PA coatings after static oxidation at 1373 K for 100 h.
~ 1.3 mg cm − 2 after 100 h of oxidation, while the mass gain for the PA coating is ~ 0.8 mg cm − 2 after the same oxidation period (Fig. 2a). This indicated that the PA coating exhibited a lower oxidation rate than the EB coating. The calculated oxidation parabolic rate constants for the EB coating and PA coating were 2.9×10−6 mg2 cm−4 s−1 and 8.0×10−7 mg2 cm−4 s−1, respectively (Fig. 2b). Compared to the oxidation of the EB coating, the scaling rate constant of the PA coating was reduced by approximately 70%. Fig. 3 shows the XRD patterns of the coating specimens after the 100 h TGA experiments. Similar oxides were detected on both coatings. α-Al2O3 was formed as the main oxide. No spinel phase but YAlOx phase was present. The surface morphologies of the oxidized coatings are shown in Fig. 4. It can clearly be seen that the alumina scales formed on the EB and PA coatings exhibited distinctly different morphologies. For the EB coating, the scale was grown with a lace-like network morphology (Fig. 4a), implying that a preferential diffusion of oxygen from oxide surface toward oxide/coating interface took place during the
3.2. TGA experiments The representative TGA curves show that the oxidation of both coatings at 1373 K followed two oxidation stages: a fast initial oxidation stage, followed by the second stage of steady-state parabolic growth (Fig. 2). The EB coating yielded a mass gain of
Fig. 2. Oxidation kinetics of the EB and PA coatings in air at 1373 K for 100 h.
Fig. 4. SEM images for the surfaces of the EB and PA coatings after oxidation at 1373 K for 100 h: (a) EB coating; (b) PA coating.
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high-temperature oxidation [23]. On the contrary, the oxidized PA coating was covered with numerous blunt needlelike oxides (Fig. 4b).
3.3. Short-term oxidation tests It has been well recognized that the TGO formed at the initial oxidation stage plays an important role on its subsequent thickening [24]. Therefore, short-term oxidation tests for different periods were performed to investigate the scale growth at the initial oxidation stage. Four time points were chosen based on the slope change of the TGA plot: after 10 min of oxidation, i.e. during the rapid increase in mass gain; after 1 h and 5 h of oxidation, i.e. during the time period in which the mass gain dropped drastically; and after 10 h, i.e. during the steady-state period of oxidation.
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Fig. 5 shows the surface morphologies of the coating specimens after different periods of oxidation. From Fig. 5a and b, it can be observed that after oxidation for 10 min at 1373 K, oxide was formed on the surface of both coatings. Lace-like oxides and whisker-like oxides were developed on the surface of the EB coating and PA coating, respectively. As the oxidation continued, the scale morphology didn't change significantly for the EB coating (Fig. 5c, e and g). On the other hand, the scale formed on the PA coating after different oxidation times exhibited an obvious morphology evolution. After oxidation for 10 min, whisker-like oxides partially covered the coating surface (Fig. 5b). With continued oxidation, the scale maintained the whisker-like morphology but grew denser and denser until 5 h (Fig. 5d and f). By 10 h oxidation, the whisker-like crystals became blunt, although the scale kept on thickening (Fig. 5h). Examination of the fractured sections revealed that the scale for the EB coating after 5 h of oxidation was composed largely of equiaxed grains
Fig. 5. SEM images for the surfaces of the EB and PA coatings after short-term oxidation tests for 10 min, 1 h, 5 h and 10 h at 1373 K.
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(Fig. 6a), whereas on the PA coating it mainly consisted of θ-Al2O3 whiskers (Fig. 6b). After 100 h of oxidation, although growing thicker (with a thickness of ~6.5 μm), the scale on the EB coating still exhibited an equiaxed structure. On the other hand, the θ-Al2O3 whiskers on the PA coating became blunt with continued oxidation, showing a twolayered structure: an ~1 μm thick outer layer consisting of equiaxed grains and a 1.2 μm thick inner layer comprising of columnar grains (Fig. 6d). The phase evolution of alumina scales formed on the EB and PA coating surface with oxidation time at 1373 K is shown in Fig. 7, where the luminescence intensity is plotted against the wavenumber. As shown in Fig. 7a, θ-Al2O3 and α-Al2O3 signals were detected in the EB coating after 10 min of oxidation. With continued oxidation, the luminescence from the θ-Al2O3 decreased in intensity and by 5 h was barely discernible, indicating almost complete transformation to the stable α-Al2O3 phase. In the case of PA coating, only θ-Al2O3 but no α-Al2O3 was detected in the very early stage (10 min). Although the subsequent phase transformation sequence was similar to that of the EB coating, the times for the transformations were significantly longer. It took ~ 10 h to accomplish the θ → α transformation compared with ~5 h for the EB coating. 4. Discussion 4.1. Effect of coating microstructure on oxidation resistance In most cases, the microstructure of the oxide scale is a decisive factor for the oxidation resistance of the coating [25–27]. During the initial oxidation stage, the fast growing θ-Al2O3 that formed on both coatings led to a rapid mass gain. However, the θ → α transformation, which was detected after 1 h of oxidation, lowered down the scaling rate. It has been reported that the initial θ → α phase transformation occurs at the oxide/metal interface [28]. Once a continuous α-Al2O3 layer is formed, the subsequent oxidation will be controlled by diffusion through the α-Al2O3 layer. This was supported by the observation of the drastic drop of the scaling rate after ~1 h of oxidation for both
coatings. For the EB coating, a faster θ → α transformation occurred, but this fast transformation was accompanied with the formation of cracks in the oxide scale, as shown in Fig. 6a (indicated by white arrow). These cracks were the result of the development of tensile stress in the oxide scale, induced by the ~14% volume reduction associated with the θ → α transformation. The presence of micro-cracks allowed for the fast growth of new oxides and consequently accelerated the oxidation rate [26]. In contrast, a slower θ→α transition could lead to an extended lifetime of the θ phase layer, which was proposed to be beneficial to the relaxation of the stresses created by the growth of α-Al2O3. During the steady-state oxidation stage, the PA coating exhibited a better oxidation resistance compared with the EB coating. Former researchers have reported that a slow-growing and adherent α-Al2O3 layer on the coating surface can be promoted by performing a preoxidation treatment under a low-oxygen partial pressure [24,27]. The explanation in the work for the better oxidation resistance after pre-oxidation was that if θ-Al2O3 was well developed in the scale, the finally obtained α-Al2O3 grains would be large and thus decrease the oxidation rate, since the scale grew predominantly by the diffusion of oxygen along the α-Al2O3 grain boundaries. A similar explanation applied to the oxidation behavior of the PA coating in the present work, except that, for the PA coating studied here, the formation of metastable θ-Al2O3 was promoted even at atmospheric air pressure. Accordingly, the lower oxidation rate for the PA coating was attributed to the formation of a more protective scale with larger α-Al2O3 grains, as compared with the scale composed of quickly transformed, fine α-Al2O3 grains. It was also noted from the XRD results (Fig. 3) that a higher amount of YAlOx was formed in the EB coating after 100 h of oxidation. Since oxygen diffuses faster through YAlOx than through alumina, the increased amount of YAlOx in the scale also contributed to the higher oxidation rate of the EB coating. This can be interpreted that the columnar grain boundaries in the EB coating provided plenty of high-diffusivity paths for the outward diffusion of Y. In such a case, Y migrated rapidly along the grain boundaries to the coating surface and thickened the scale. Within the PA coating, the equiaxed grain
Fig. 6. SEM images for the fractured sections of the EB and PA coatings after short-term oxidation tests for 5 h and TGA experiments for 100 h at 1373 K. (Note that the images are shown at different magnifications.)
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Fig. 7. Evolution of luminescence spectrum for the EB and PA coatings after short-term oxidation tests at 1373 K: (a) EB coating; (b) PA coating.
reduced the outward diffusion rate of Y in some degree, then less YAlOx was formed. Similar results and discussions have been reported in Ref. [14].
4.2. Effect of coating microstructure on TGO formation The explanation given above raises the following question: why was the formation of θ-Al2O3 greatly enhanced in the PA coating? It is known that certain alloying elements can affect the kinetics of the θ→α transformation (for instance, Cr accelerates the θ→α transformation and Y retards the θ→α transformation) [25]. However, it can be noted that the EB and PA coatings had similar average compositions, thus the element competition effect can be eliminated. Therefore, it can be inferred that it was the coating microstructure that played the decisive role in the oxidation behavior. Another retarded θ → α transformation was observed at local regions of a plasma sprayed NiCoCrAlY coating [29]. A conclusion was drawn that α-Al2O3 and θ-Al2O3 initially formed on different phases in the coating and that there was then a slow transformation during oxidation of the θ-Al2O3 phase to α-Al2O3 correlated to the β-NiAl phase. Contrary to this, a homogeneous θ-Al2O3 scale was developed initially on the PA coating. The micro-lamella β phase distributed in the γ matrix within the PA coating was more beneficial in promoting the outward diffusion of Al to the coating surface. This is attributed to the following two reasons: (1) the larger volume percentage and smaller size of the β phase result in a higher density of γ/β phase boundaries in the coating, thereby providing fast diffusion paths for Al toward the coating surface [12,30], and (2) the nucleation of θ-Al2O3 is favored on the β phase of the coating surface due to preferred (epitaxial) orientation relationships between θ-Al2O3 and β-NiAl [31]. The finely spaced γ/β phase promoted the homogeneous growth of θ-Al2O3.
5. Conclusions A novel structured NiCoCrAlY coating was produced by plasma activated electron beam-physical vapor deposition (PA EB-PVD). Within the coating, the well-arranged β-NiAl phase precipitated as small lamellae (less than 1 μm in thickness) in the γ matrix, predominantly
oriented perpendicular to the coating surface. The conclusions drawn from the oxidation investigations are given as follows: (1) The scaling rate constant (8.0 × 10 − 7 mg 2 cm − 4 s − 1) during the steady-state of the novel coating was reduced by approximately 70% as compared with the conventional coating (2.9 × 10 − 6 mg 2 cm − 4 s − 1) of similar composition produced by electron beam-physical vapor deposition (EB-PVD). (2) The vertical arrangement of the micro-lamellae in the PA coating that provided fast diffusion paths promoted the outward diffusion of Al cations and the subsequent formation of well developed θ-Al2O3 scales. (3) The α-Al2O3 that transformed from the θ-Al2O3 grains grew at a much slower rate during the steady-state oxidation than the scale developed on a conventional NiCoCrAlY coating deposited by EB-PVD. Acknowledgments This work was financially supported by the National Basic Research Program (973 Program) of China under grant No. 2010CB631200 and the National Nature Science Foundations of China (NSFC, No.50771009, No. 50731001 and No. 51071013). References [1] P. Richer, M. Yandouzi, L. Beauvais, B. Jodoin, Surf. Coat. Technol. 204 (2010) 3962. [2] Y. Li, C.J. Li, G.J. Yang, L.K. Xing, Surf. Coat. Technol. 205 (2010) 2225. [3] S.M. Jiang, C.Z. Xu, H.Q. Li, J. Ma, J. Gong, C. Sun, Corros. Sci. 52 (2011) 1746. [4] D. Naumenko, V. Shemet, L. Singheiser, W.J. Quadakkers, J. Mater. Sci. 44 (2009) 1687. [5] J. Ma, S.M. Jiang, H.Q. Li, W.X. Wang, J. Gong, C. Sun, Corros. Sci. 53 (2011) 1417. [6] L.C. Chen, C. Zhang, Z.G. Yang, Corros. Sci. 53 (2011) 374. [7] B. Gleeson, N. Mu, S. Hayashi, J. Mater. Sci. 44 (2009) 1704. [8] H. Guo, Y. Cui, H. Peng, S. Gong, Corros. Sci. 52 (2011) 1440. [9] H. Lan, Z.G. Yang, Z.X. Xia, Y.D. Zhang, C. Zhang, Corros. Sci. 53 (2011) 1476. [10] J.F. Loersch, J.W. Neal, U. S. Pat. 4514469, (1981). [11] R.G. Wellman, A. Scrivani, G. Rizzi, A. Weisenburger, F.H. Tenailleau, J.R. Nicholls, Surf. Coat. Technol. 202 (2007) 709. [12] T.J. Nijdam, C. Kwakernaak, W.G. Sloof, Metall. Mater. Trans. A 37 (2006) 683. [13] Z. Liu, W. Gao, K. Dahm, F. Wang, Scr. Mater. 37 (1997) 1551. [14] H. Peng, H. Guo, J. He, S. Gong, Surf. Coat. Technol. 205 (2011) 4658. [15] H.A. Beale, T.E. Strangman, E.W. Taylor, U. S. Pat. 4109061, (1978).
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