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High temperature corrosion behavior of a multilayer CrAlN coating prepared by magnetron sputtering method on a K38G alloy
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Surface & Coatings Technology 202 (2008) 1985 – 1993 www.elsevier.com/locate/surfcoat
High temperature corrosion behavior of a multilayer CrAlN coating prepared by magnetron sputtering method on a K38G alloy Tianpeng Li a,b , Yanchun Zhou a,⁎, Meishuan Li a , Zhongping Li c a
High-performance Ceramic Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China b Graduate School of Chinese Academy of Sciences, Beijing, 100039, China c National Key Laboratory of Advanced Functional Composite Materials, Beijing, 100076, China Received 3 June 2007; accepted in revised form 20 August 2007 Available online 25 August 2007
Abstract A multilayer CrAlN coating of Cr0.58Al0.42N/Cr0.84Al0.16N/Cr0.51Al0.49N has been fabricated by a reactive magnetron sputtering method. It consists of a bonding layer, a Cr-rich intermediate layer and an Al-rich outer layer. The multilayer structure provides the coating with good protection against different types of high temperature corrosion, i.e., high temperature oxidation and hot corrosion. The outer Al-rich layer gives the coating good oxidation resistance at 1000 and 1100 °C due to the formation of a continuous alumina scale. The parabolic rate constants of the coated samples decrease by about 2 orders of magnitude compared with that of the bare alloy samples. The intermediate Cr-rich layer can form a Cr2O3 scale to provide good protection under the hot corrosion conditions in the Na2SO4 salt fluxing at 900, 950 and 1000 °C. The incubation period of the hot corrosion extends several times longer when the alloy was coated by the multilayer coating at the three selected temperatures. © 2007 Elsevier B.V. All rights reserved. Keywords: Multilayer coating; Oxidation; Hot corrosion; Magnetron sputtering
1. Introduction Since 1965, the compositions of nickel-based alloys have changed greatly. The content of refractory elements increased from 8 to 20 wt.%, but the content of chromium was reduced drastically from about 15 wt.% to about 3 wt.% and the aluminum content has change only a little [1]. All of these developments have been undertaken to improve mechanical properties of the alloys at higher temperatures. However, the low chromium and aluminum contents mean that these materials do not have the necessary intrinsic resistance to oxidation and corrosion required for the long term operation associated with gas turbine applications. A number of high-performance coatings, such as diffusion coatings [2,3], modified diffusion coatings [4–8], MCrAlY overlay coatings [9,10], thermal barrier coatings [11–15] and smart ⁎ Corresponding author. Tel.: +86 24 23971765; fax: +86 24 23891320. E-mail address: [email protected] (Y. Zhou). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.08.025
overlay coatings [16,17], have been developed. Chromium and aluminum are two main elements for the formation of protective scales in the coatings. Chromia can volatilize as CrO3 compound at temperatures higher than about 1000 °C. An aluminum addition for oxidation resistance is preferred at and above 1000 °C for key components. But the hot corrosion resistance of Al2O3 against molten salts is not as good as Cr2O3 [18–21]. In this work, a multilayer CrAlN coating has been fabricated using a reactive magnetron sputtering method. The multilayer coating consists of a bonding layer (Cr0.58Al042N) Table 1 Nominal composition of a K38G alloy used as substrate in the present work Element
C
Cr
Co
W
Mo
Al
Content (wt. %) 0.13–0.20 15.3–16.8 8.0–9.0 2.3–2.9 1.4–2.0 3.5–4.5 Element
C
Content (wt. %) 3.2–4.0
Cr
Co
b0.20
0.4–1.0 1.4–2.0
W
Mo
Al
0.05–0.15 balance
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Table 2 Deposition parameters for the multilayer coating Parameter
Value
Size of Cr and Al targets Size of substrate Substrate temperature Gas flow Work pressure Base pressure Distance from each target to the substrate Power density of the Al targets for the bonded layer Power density of the Cr targets for the bonded layer Power density of the Al targets for the Cr-rich layer Power density of the Cr targets for the Cr-rich layer Power density of the Al targets for the Al-rich layer Power density of the Cr targets for the Al-rich layer
Φ60 mm 10 mm × 10 mm × 2 mm 300 °C Ar: N2 = 8: 12 (SCCM) 0.3 Pa 2 × 10− 4 Pa 70 mm 4.2 W/cm2 4.2 W/cm2 2.1 W/cm2 4.2 W/cm2 4.2 W/cm2 3.6 W/cm2
adjacent to the substrate, a Cr-rich intermediate layer (Cr0.84Al0.16N) and an Al-rich outer layer (Cr0.51Al0.49N) with the thickness of ∼ 2 μm, ∼ 7 μm and ∼ 5 μm, respectively. It is expected that the bonding layer can improve the mismatch of the thermal expansion coefficient between the Crrich layer and the substrate because of the higher thermal expansion coefficient of AlN than that of CrN. The multilayer coating is superior against hot corrosion in molten salts to that of the single-layer CrAlN coating [22,23]. In comparison with the single-layer CrAlN coating, the multilayer coating demonstrates a more pseudo-intelligent response to the corrosion conditions, i.e., Cr2O3 and Al2O3 scales form in the conditions of hot corrosion and elevated temperature oxidation, respectively. Therefore, the coating can be used in a wide temperature range and would be ideal for turbine engine applications. This multilayer CrAlN coating can be regarded as another type of novel smart coatings. 2. Experimental procedure A K38G alloy was chosen as the substrate in the present work. The composition of the alloy is listed in Table 1. The specimens
Fig. 2. Surface morphology of the multilayer coating (a) and cross-section microstructure of the coating together with the EDS line profiles of the elements Al, Cr, and Ni for the as-deposited coating (b).
Fig. 1. XRD pattern from the surface of the multilayer coating.
were cut into slices of the dimensions 10 mm × 10 mm × 2 mm, and their surfaces were ground down to 1200 grit SiC paper and then polished with a 0.5 μm diamond paste. Then the samples were degreased ultrasonically in acetone, cleaned by ethanol and dried in air. Cr–Al–N coatings were deposited in a JGP560C14 magnetron sputtering system (SKY Technology Development Co. Ltd, Shenyang, China) by reactive magnetron co-sputtering from two targets of chromium (purity 99.99%) and aluminum (purity 99.99%) at 300 °C in a D.C mode. The flow rates of the
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Table 3 Parabolic rate constants of the coated samples oxidized at 1000–1100 °C and that of the bare K38G alloy at 1000–1100 °C for 20 h in air Sample
Coated sample K38G alloy
Temperature 1000 °C 1050 °C kp (×10− 11 kg2 m− 4s− 1)
1100 °C
3.08 108
4.05 149
3.58 135
reactive gas (N2, 99.999%) and the inert gas (Ar, 99.999%) were controlled by independent mass-flow controllers, and the gases were mixed before they entered the deposition chamber. The substrates were immobilized during the deposition and the deposition parameters of the multilayer coating are listed in Table 2. An X-ray diffractometer (XRD, Rigaku D/max-2400, Japan) was used to identify the phase composition of the as-deposited coating and the oxide scales. A scanning electron microscope (SEM, LEO Super 35, Germany) equipped with energy dispersive X-ray spectroscopy (EDS, Oxford INCA x-sight, England) was used to investigate the surface morphologies and
Fig. 4. XRD patterns of the coated samples after oxidation at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C.
the cross-sections of the coated specimens before and after oxidation or hot corrosion at different temperatures. Rietveld method [24] was used to calculate the phase contents of the as-
Fig. 3. Squares of the mass gain as a function of oxidation time for the coated and bare substrate at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C during 20 hours; (d) is the kinetics of the coated and uncoated samples oxidized at 1000 °C during 300 hours.
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deposited coating, which was accomplished in a DBWS program in the Cerius2 computational program for material research (Molecular Simulation Inc., USA). The isothermal-oxidation and hot corrosion tests were performed with use of a Setsys 16/18 thermal balance (Setaram, France) at 1000–1100 °C and 900–1000 °C in air for 20 h, respectively. The samples were heated to the desired temperatures at the rate of 30 °C/min and then the mass changes were recorded as a function of time. The samples were coated by Na2SO4 with 3 ± 0.2 mg/cm 2 for hot corrosion tests. To investigate the long term oxidation behaviors, an oxidation test with the duration of 300 h was conduced at 1000 °C in static air. An analytical balance with a sensitivity of 10− 5 g was used to measure the weight of the specimens during the test.
3. Results 3.1. Characterization of the as-deposited coating Fig. 1 gives the XRD pattern of the as-deposited coating. The coating consists mainly of rock-salt type B1 structure with a small admixture of AlN phase of wurtzite type B4 structure phase. The quantitative analysis of the XRD pattern confirmed that the content of B4 h-AlN phase was about 5 wt.%. The crystallites of the CrN phase demonstrate a preferred orientation of {111} planes parallel to the coating surface. The surface morphology of the coating, its cross-sectional microstructure and EDS analyses of the as-deposited coating are presented in Fig. 2(a) and (b), respectively. The coating is compact with no visible crack. It is clearly seen that the coating consists of three
Fig. 5. Surface morphology of the coated samples after oxidation at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C. (d) and (e) are EDS spectra of disk-like and irregular shap grains, respectively.
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oxide scale of the coating is very thin, therefore only the peaks corresponding to the coating are observed. However, at 1050 °C and 1100 °C, α-Al2O3 phase and a little amount of Cr2O3 is dominating in the XRD patterns. The SEM images of the surface morphologies of the coated samples after oxidation at 1000 °C, 1050 °C and 1100 °C are presented in Fig. 5(a), (b) and (c), respectively. After oxidation
Fig. 6. SEM image of the cross-section of the coated sample after oxidation at 1050 °C together with the distribution of the elements Al, Cr, Ni, Ti and O.
layers. From the semiquantitative EDS analysis of Cr to Al ratio, the compositions of the three layers are calculated as follows: Cr0.6Al0.4N for the bonding layer, Cr0.8Al0.2N for the intermediate one and Cr0.5Al0.5N for the superficial layer. We must acknowledge that no effort was made to determine the N content. The corresponding thicknesses of the layers are ∼ 2 μm, ∼ 7 μm and ∼5 μm, respectively (see Fig. 2b). 3.2. Isothermal oxidation The squares of the mass gain as a function of oxidation time for the coated and bare alloys at 1000 °C, 1050 °C and 1100 °C are depicted in Fig. 3(a), (b) and (c), respectively. The weight gains of the coated alloy are much less than that of the bare alloy during the oxidation at the three temperatures. The straight line relationship indicates that the oxidation kinetics of the coated and bare alloy obey the parabolic rate law roughly. The parabolic rate constants (kp) have been determined and are listed in Table 3. It is apparent that for the alloy coated with the multilayer coating, the parabolic rate constants for oxidation decrease by about 2 orders of magnitude at the three temperatures. The kinetics of oxidation at 1000 °C for coated and uncoated samples during 300 h is presented in Fig. 3 (d). The mass gain of the coated sample is much less than that of the uncoated one, which approves that the coating has a good long-term resistance to oxidation. It can be clearly seen that the uncoated sample experiences a weight loss in the last ∼150 h, which indicated the breakdown of the scale. The XRD patterns of the samples after oxidation at 1000 °C, 1050 °C and 1100 °C are presented in Fig. 4. At 1000 °C, the
Fig. 7. Hot corrosion kinetics of the bare and coated alloy at (a) 900 °C, (b) 950 °C and (c) 1000 °C.
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Fig. 8. XRD patterns of the coated samples after hot corrosion in Na2SO4 salt flux at 900 °C, 950 °C and 1000 °C.
at 1000 °C, the oxide grains with disk-like shape as well as those with irregular shape appear. The grains are likely Cr2O3 and Al2O3, respectively, based on the analysis of EDS spectra shown in Fig. 5 (d) and (e). When the temperature increases to 1050 °C, the amount of irregular shape Al2O3 grains increases significantly and the disk-like Cr2O3 grains almost disappear. At 1100 °C, the size of Al2O3 grains increases slightly. The cross-section microstructure of the sample oxidized at 1050 °C is depicted in Fig. 6 together with the distribution of the elements Al, Cr, Ni, Ti and O. It can be clearly seen that an aluminum-and oxygen-rich layer, which may be Al2O3, formed on the surface of the coating and an interdiffusion between the coating and the substrate is not observed. A titanium-enriched zone is identified in the inner bonding layer of the coating, which indicates that the intermediate chromium enriched layer limits effectively the outward diffusion of Ti atoms. An oxygen rich layer forms at the inner bonding layer, which was resulted from the oxygen diffusion via the defects, such as crack, in the intermediate Cr rich layer. However, the oxygen was not observed to transport into the substrate alloy, they are trapped by the aluminum atoms in the inner bonding layer. 3.3. Hot corrosion The hot corrosion kinetics of the coated samples and the bare K38G alloy at 900 °C, 950 °C and 1000 °C are presented in Fig. 7(a), (b) and (c), respectively. When the bare alloy is exposed to oxidation at 900 °C and 950 °C, two periods can be identified from the hot corrosion kinetics curves. These two periods have been named an incubation period and a propagation period, respectively. The hot corrosion kinetics of the bare alloy exposed at 1000 °C can be divided into three periods, which are labeled as incubation, propagation and stabilization periods. However, all the coated samples exposed at 900 °C, 950 °C and 1000 °C exhibit only the incubation period during the whole exposed time of 20 hours. The kinetics curves of the coated samples indicate that the incubation period of the coated samples in the molten Na2SO4 flux is longer than 20 h at the three selected temperatures.
The XRD patterns of the coated samples after hot corrosion in Na2SO4 salt fluxing at 900 °C, 950 °C and 1000 °C are shown in Fig. 8. The phase compositions of the scales are mainly Cr2O3 with a trace of α-Al2O3. Fig. 9(a) and (b) reveal the surface morphology of the coated samples after hot corrosion at 900 °C and 1000 °C, respectively. A two-layer microstructure of the corrosion products can be distinguished on the cross-section of the samples, i.e., a loose layer on the top of the samples and a compact layer beneath it. The loose layer on the top of the sample is partly bonded with the underneath compact layer. The enlarged images of the compact bottom layer and of the loose top layer after hot corrosion tests at 1000 °C are given in Fig. 9 (c) and (d), respectively. The EDS spectra of Fig. 9 (c) and (d) are shown in Fig. 9 (e) and (f), respectively. The phase compositions of the loose layer and the compact layer were determined as α-Al2O3 and Cr2O3, respectively, based on the analysis of EDS and XRD. The cross-section microstructure as well as the distribution of the elements Al, Ti, Cr, O, Ni and S of the bare and coated alloy after exposition to hot corrosion in fused Na2SO4 at 1000 °C are shown in Fig. 10(a) and (b), respectively. One can see that the outer layer enriched in Al has been dissolved and almost completely consumed by the molten Na2SO4 salt, whereas the intermediate layer rich in Cr atoms is compact and provides a good protection against the molten salt. In comparison with the coated sample the atoms of sulfur (S) penetrate into the bare alloy more deeply, and the grain boundaries and other defects in the substrate alloy seem to act as short-circuits for transport of the sulfur atoms. 4. Discussion 4.1. Isothermal oxidation processes The outer layer of the multilayer coating has a high aluminum content, which was designed to resist the high temperature oxidation. Similar to the single-layer CrAlN coatings with a high aluminum content investigated in a previous work [22], an alumina-enriched scale can form at elevated temperatures due to the high aluminum activity (content) in the coating and the greater negative value of the Gibbs energy for the formation of alumina than that for chromia. In other words, an alumina-rich scale formation is preferential both from the kinetics as well as from the thermodynamics point of view. This is why the coating gives a good protection against oxidation at elevated temperatures. The chromium-rich intermediate layer behaves as a diffusion barrier to limit the outward diffusion of the atoms of the substrate alloy elements to the outer Al-rich layer. It is also believed that the chromium-rich interlayer can limit the loss of aluminum by diffusing into the substrate. No apparent interdiffusion is observed between the coating and the substrate, as shown in Fig. 6. Therefore, the aluminum depletion due to the interdiffusion at the interface between the coating and the substrate is minimized. As a result the high alumina content in the scale and its beneficial anti-oxidizing properties can be maintained for a long time. It is expected that the lifetime of the coated components in gas turbine engines would be enhanced.
T. Li et al. / Surface & Coatings Technology 202 (2008) 1985–1993
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Fig. 9. Surface morphology of the coated samples after hot corrosion at (a) 900 °C, (b) 1000 °C. (c) and (d) are enlarged surface images of the compact bottom layer and the loose top layer after hot corrosion test at 1000 °C, respectively. (e) and (f) are the EDS spectra of (c) and (d), respectively.
4.2. Hot corrosion processes At temperatures below 1000 °C, the components can suffer another type of aggressive attack from the environment, the hot corrosion. During the incubation period, the mass gain is
relatively low and, the oxides on the surface do not break down. Once the oxide scale breaks down and its self-healing is no longer possible, the propagation period begins and a rapid deterioration of the coating takes place. Due to a lower resistance of alumina to the salt fluxing in comparison to that of
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Fig. 10. Cross-sectional microstructure of the bare (a) and coated sample (b) together with the distribution of elements Al, Ti, Cr, O, Ni and S after exposure in fused Na2SO4 at 1000 °C.
chromia, the outer Al-rich layer of the multilayer coating does not protect against the salt fluxing effectively and will dissolve in the molten salts, and precipitate as loose oxides at the interface between the salt flux and the gas environment. As a result the outer Al-rich layer is consumed rapidly and the intermediate Cr-rich layer exposes to the salt fluxing. The intermediate chromium-rich layer can readily form a continuous chromia scale and provide a rapid crack-healing route preventing the breakdown of scale. Since the coating provides protective surface oxide scales with the crack-healing properties, the incubation period can be extended greatly and the initiation of the propagation period can be increased to a great extent. This mechanism is crucial for the design of the long
lifetime components. As shown in Fig. 7, the propagation periods were not observed for all the coated samples through the exposed time in the present work, which indicates that the incubation periods for the coated samples are extended over 20 h. In comparison with the bare substrate, the incubation period of the coated alloy increases several times. For the samples exposed at 1000 °C, the incubation period of the coated sample is more than 10 times longer than that of the bare substrate. Moreover, from Fig. 7 (a), (b) and (c) it can be clearly seen that the coated samples demonstrate a slight mass loss, which was resulted from the volatilization of chromia. Erdös et al. [25] observed the mass loss during the hot corrosion of IN 939 in
T. Li et al. / Surface & Coatings Technology 202 (2008) 1985–1993
molten (Na0.9, K0.1)2SO4 at 900 °C. They assumed that the volatilization took place by the formation and subsequent decomposition of sodiumdichromate to chromia. In the present work, N2 emitting also contributed to the mass loss. From the cross-section microstructure and from the analysis of the distribution of the elements (as shown in Fig. 10(a) and (b)), the penetration of sulfur (S) is very deep for the bare substrate. In contrast, no apparent penetration of sulfur (S) was observed for the coated substrate, which indicates that the multilayer coating prevents effectively the inward diffusion of S atoms into the substrate. The results in the present paper demonstrate that the multilayer coating can respond in a pseudo-intelligent mode to the aggressive attackes of the environment, i.e., to oxidation at elevated temperatures as well as to hot corrosion at intermediate temperatures. Therefore, the multilayer CrAlN coating can be regarded as a novel type of smart coatings. 5. Conclusion A multilayer CrAlN coating of Cr0.58Al0.42N/Cr0.84Al0.16N/ Cr0.51Al0.49N was synthesized by a magnetron sputtering method on a K38G alloy. The isothermal oxidation at 1000 °C, 1050 °C and 1100 °C and the hot corrosion in a Na2SO4 salt fluxing at 900 °C, 950 °C and 1000 °C are investigated in air. Zhu [23] has investigated the corrosion behaviors of CrAlN coatings at similar temperatures. The present work demonstrated comparative resistance to oxidation but much more superior resistance to hot corrosion. The following conclusions can be drawn from the investigations performed in the present work: (1) The resistance to isothermal oxidation of the nickel-based K38G substrate alloy is enhanced significantly by the multilayer coating. The parabolic rate constants for the coated samples decrease by 2 orders of magnitude at the temperatures of 1000 °C and 1100 °C. The outer Al-rich layer contributes to the good resistance to isothermal oxidation due to formation of a continuous alumina scale on the coating's surface. (2) The intermediate Cr-rich layer can provide a diffusion barrier for the outward diffusion of the alloy elements from the substrate to the coating and the inward diffusion of Al from the coating to the substrate. (3) The resistance to hot corrosion is improved when the alloy is coated by the multilayer coating. The incubation
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period is extended from approximately 3 hours for the uncoated alloy to over 20 hours for the coated samples. The intermediate Cr-rich layer is responsible for the good resistance to the hot corrosion in molten Na2SO4 salt flux as a continuous chromia scale forms on the surface and provides a rapid route for healing the breakdown of the scale. Acknowledgements This work was supported by the National Outstanding Young Scientist Foundation of China for Y. C. Zhou under grant No. 59925208, Natural Science Foundation of China under grant No. 50232040, 50571106, 50072034, 90403027, 863 project, and High-tech Bureau of the Chinese Academy of Sciences. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]
M.J. Pomeroy, Mater. Des. 26 (2005) 223. A.R. Castle, D.R. Gabe, Int. Mater. Rev. 44 (1999) 37. A. Elsawy, H. Soda, Werkst. Korros. 56 (2005) 669. G.R. Krishna, D.K. Das, V. Singh, S.V. Joshi, Mater. Sci. Eng., A 251 (1998) 40. C. Leyens, B.A. Pint, I.G. Wright, Surf. Coat. Technol. 133 (2000) 15. Y. Zhang, J.A. Haynes, W.Y. Lee, I.G. Wright, B.A. Pint, Metall. Mater. Trans., A 30 (1999) 2679. Y. Zhang, J.A. Haynes, W.Y. Lee, I.G. Wright, B.A. Pint, Metall. Mater. Trans., A 32 (2001) 1927. A.L. Purvis, B.M. Warnes, Surf. Coat. Technol. 146 (2001) 1. G.W. Goward, Surf. Coat. Technol. 109 (1998) 73. Y. Itoh, M. Tamura, J. Eng. Gas Turbine Power 121 (1999) 476. W.J. Brindley, R.A. Miller, Adv. Mater. Process. 8 (1989) 29. Y.H. Sohn, R.R. Biederman, J.R. Sisson, Thin Solid Films 250 (1994) 1. E. Shillington, D.R. Clarke, Acta Mater. 47 (1999) 1297. A.G. Evans, D.R. Mumm, J.W. Meiter, F.S. Pettit, Prog. Mater. Sci. 46 (2001) 505. V.K. Tolpygo, D.R. Clarke, Acta Mater. 48 (2000) 3283. J.R. Nicholls, Mater. World 4 (1996) 19. J.R. Nicholls, N.J. Simms, W. Chan, H.E. Evans, Surf. Coat. Technol. 149 (2002) 236. Y.S. Zhang, R.A. Rapp, Corrosion 43 (1987) 348. Y.S. Zhang, R.A. Rapp, J. Electrochem. Soc. 132 (1985) 2498. Y.S. Zhang, J. Electrochem. Soc. 133 (1986) 655. R.A. Rapp, Y.S. Zhang, JOM 46 (1994) 47. T.P. Li, X.H. Yin, M.S. Li, Y.C. Zhou, Oxid. Met. 68 (2007) 193. M. Zhu, M. Li, Y. Zhou, Surf. Coat. Technol. 201 (2006) 2878. R.A. Young, The Rietveld Method, Oxford University Press, Oxford, U.K., 1993, p. 1. E. Erdös, H. Altorter, E. Denzler, Werkst. Korros. 33 (1982) 373.
Surface & Coatings Technology 202 (2008) 1985 – 1993 www.elsevier.com/locate/surfcoat
High temperature corrosion behavior of a multilayer CrAlN coating prepared by magnetron sputtering method on a K38G alloy Tianpeng Li a,b , Yanchun Zhou a,⁎, Meishuan Li a , Zhongping Li c a
High-performance Ceramic Division, Shenyang National Laboratory for Materials Science, Institute of Metal Research, Chinese Academy of Sciences, 72 Wenhua Road, Shenyang, 110016, China b Graduate School of Chinese Academy of Sciences, Beijing, 100039, China c National Key Laboratory of Advanced Functional Composite Materials, Beijing, 100076, China Received 3 June 2007; accepted in revised form 20 August 2007 Available online 25 August 2007
Abstract A multilayer CrAlN coating of Cr0.58Al0.42N/Cr0.84Al0.16N/Cr0.51Al0.49N has been fabricated by a reactive magnetron sputtering method. It consists of a bonding layer, a Cr-rich intermediate layer and an Al-rich outer layer. The multilayer structure provides the coating with good protection against different types of high temperature corrosion, i.e., high temperature oxidation and hot corrosion. The outer Al-rich layer gives the coating good oxidation resistance at 1000 and 1100 °C due to the formation of a continuous alumina scale. The parabolic rate constants of the coated samples decrease by about 2 orders of magnitude compared with that of the bare alloy samples. The intermediate Cr-rich layer can form a Cr2O3 scale to provide good protection under the hot corrosion conditions in the Na2SO4 salt fluxing at 900, 950 and 1000 °C. The incubation period of the hot corrosion extends several times longer when the alloy was coated by the multilayer coating at the three selected temperatures. © 2007 Elsevier B.V. All rights reserved. Keywords: Multilayer coating; Oxidation; Hot corrosion; Magnetron sputtering
1. Introduction Since 1965, the compositions of nickel-based alloys have changed greatly. The content of refractory elements increased from 8 to 20 wt.%, but the content of chromium was reduced drastically from about 15 wt.% to about 3 wt.% and the aluminum content has change only a little [1]. All of these developments have been undertaken to improve mechanical properties of the alloys at higher temperatures. However, the low chromium and aluminum contents mean that these materials do not have the necessary intrinsic resistance to oxidation and corrosion required for the long term operation associated with gas turbine applications. A number of high-performance coatings, such as diffusion coatings [2,3], modified diffusion coatings [4–8], MCrAlY overlay coatings [9,10], thermal barrier coatings [11–15] and smart ⁎ Corresponding author. Tel.: +86 24 23971765; fax: +86 24 23891320. E-mail address: [email protected] (Y. Zhou). 0257-8972/$ - see front matter © 2007 Elsevier B.V. All rights reserved. doi:10.1016/j.surfcoat.2007.08.025
overlay coatings [16,17], have been developed. Chromium and aluminum are two main elements for the formation of protective scales in the coatings. Chromia can volatilize as CrO3 compound at temperatures higher than about 1000 °C. An aluminum addition for oxidation resistance is preferred at and above 1000 °C for key components. But the hot corrosion resistance of Al2O3 against molten salts is not as good as Cr2O3 [18–21]. In this work, a multilayer CrAlN coating has been fabricated using a reactive magnetron sputtering method. The multilayer coating consists of a bonding layer (Cr0.58Al042N) Table 1 Nominal composition of a K38G alloy used as substrate in the present work Element
C
Cr
Co
W
Mo
Al
Content (wt. %) 0.13–0.20 15.3–16.8 8.0–9.0 2.3–2.9 1.4–2.0 3.5–4.5 Element
C
Content (wt. %) 3.2–4.0
Cr
Co
b0.20
0.4–1.0 1.4–2.0
W
Mo
Al
0.05–0.15 balance
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T. Li et al. / Surface & Coatings Technology 202 (2008) 1985–1993
Table 2 Deposition parameters for the multilayer coating Parameter
Value
Size of Cr and Al targets Size of substrate Substrate temperature Gas flow Work pressure Base pressure Distance from each target to the substrate Power density of the Al targets for the bonded layer Power density of the Cr targets for the bonded layer Power density of the Al targets for the Cr-rich layer Power density of the Cr targets for the Cr-rich layer Power density of the Al targets for the Al-rich layer Power density of the Cr targets for the Al-rich layer
Φ60 mm 10 mm × 10 mm × 2 mm 300 °C Ar: N2 = 8: 12 (SCCM) 0.3 Pa 2 × 10− 4 Pa 70 mm 4.2 W/cm2 4.2 W/cm2 2.1 W/cm2 4.2 W/cm2 4.2 W/cm2 3.6 W/cm2
adjacent to the substrate, a Cr-rich intermediate layer (Cr0.84Al0.16N) and an Al-rich outer layer (Cr0.51Al0.49N) with the thickness of ∼ 2 μm, ∼ 7 μm and ∼ 5 μm, respectively. It is expected that the bonding layer can improve the mismatch of the thermal expansion coefficient between the Crrich layer and the substrate because of the higher thermal expansion coefficient of AlN than that of CrN. The multilayer coating is superior against hot corrosion in molten salts to that of the single-layer CrAlN coating [22,23]. In comparison with the single-layer CrAlN coating, the multilayer coating demonstrates a more pseudo-intelligent response to the corrosion conditions, i.e., Cr2O3 and Al2O3 scales form in the conditions of hot corrosion and elevated temperature oxidation, respectively. Therefore, the coating can be used in a wide temperature range and would be ideal for turbine engine applications. This multilayer CrAlN coating can be regarded as another type of novel smart coatings. 2. Experimental procedure A K38G alloy was chosen as the substrate in the present work. The composition of the alloy is listed in Table 1. The specimens
Fig. 2. Surface morphology of the multilayer coating (a) and cross-section microstructure of the coating together with the EDS line profiles of the elements Al, Cr, and Ni for the as-deposited coating (b).
Fig. 1. XRD pattern from the surface of the multilayer coating.
were cut into slices of the dimensions 10 mm × 10 mm × 2 mm, and their surfaces were ground down to 1200 grit SiC paper and then polished with a 0.5 μm diamond paste. Then the samples were degreased ultrasonically in acetone, cleaned by ethanol and dried in air. Cr–Al–N coatings were deposited in a JGP560C14 magnetron sputtering system (SKY Technology Development Co. Ltd, Shenyang, China) by reactive magnetron co-sputtering from two targets of chromium (purity 99.99%) and aluminum (purity 99.99%) at 300 °C in a D.C mode. The flow rates of the
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Table 3 Parabolic rate constants of the coated samples oxidized at 1000–1100 °C and that of the bare K38G alloy at 1000–1100 °C for 20 h in air Sample
Coated sample K38G alloy
Temperature 1000 °C 1050 °C kp (×10− 11 kg2 m− 4s− 1)
1100 °C
3.08 108
4.05 149
3.58 135
reactive gas (N2, 99.999%) and the inert gas (Ar, 99.999%) were controlled by independent mass-flow controllers, and the gases were mixed before they entered the deposition chamber. The substrates were immobilized during the deposition and the deposition parameters of the multilayer coating are listed in Table 2. An X-ray diffractometer (XRD, Rigaku D/max-2400, Japan) was used to identify the phase composition of the as-deposited coating and the oxide scales. A scanning electron microscope (SEM, LEO Super 35, Germany) equipped with energy dispersive X-ray spectroscopy (EDS, Oxford INCA x-sight, England) was used to investigate the surface morphologies and
Fig. 4. XRD patterns of the coated samples after oxidation at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C.
the cross-sections of the coated specimens before and after oxidation or hot corrosion at different temperatures. Rietveld method [24] was used to calculate the phase contents of the as-
Fig. 3. Squares of the mass gain as a function of oxidation time for the coated and bare substrate at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C during 20 hours; (d) is the kinetics of the coated and uncoated samples oxidized at 1000 °C during 300 hours.
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deposited coating, which was accomplished in a DBWS program in the Cerius2 computational program for material research (Molecular Simulation Inc., USA). The isothermal-oxidation and hot corrosion tests were performed with use of a Setsys 16/18 thermal balance (Setaram, France) at 1000–1100 °C and 900–1000 °C in air for 20 h, respectively. The samples were heated to the desired temperatures at the rate of 30 °C/min and then the mass changes were recorded as a function of time. The samples were coated by Na2SO4 with 3 ± 0.2 mg/cm 2 for hot corrosion tests. To investigate the long term oxidation behaviors, an oxidation test with the duration of 300 h was conduced at 1000 °C in static air. An analytical balance with a sensitivity of 10− 5 g was used to measure the weight of the specimens during the test.
3. Results 3.1. Characterization of the as-deposited coating Fig. 1 gives the XRD pattern of the as-deposited coating. The coating consists mainly of rock-salt type B1 structure with a small admixture of AlN phase of wurtzite type B4 structure phase. The quantitative analysis of the XRD pattern confirmed that the content of B4 h-AlN phase was about 5 wt.%. The crystallites of the CrN phase demonstrate a preferred orientation of {111} planes parallel to the coating surface. The surface morphology of the coating, its cross-sectional microstructure and EDS analyses of the as-deposited coating are presented in Fig. 2(a) and (b), respectively. The coating is compact with no visible crack. It is clearly seen that the coating consists of three
Fig. 5. Surface morphology of the coated samples after oxidation at (a) 1000 °C, (b) 1050 °C and (c) 1100 °C. (d) and (e) are EDS spectra of disk-like and irregular shap grains, respectively.
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oxide scale of the coating is very thin, therefore only the peaks corresponding to the coating are observed. However, at 1050 °C and 1100 °C, α-Al2O3 phase and a little amount of Cr2O3 is dominating in the XRD patterns. The SEM images of the surface morphologies of the coated samples after oxidation at 1000 °C, 1050 °C and 1100 °C are presented in Fig. 5(a), (b) and (c), respectively. After oxidation
Fig. 6. SEM image of the cross-section of the coated sample after oxidation at 1050 °C together with the distribution of the elements Al, Cr, Ni, Ti and O.
layers. From the semiquantitative EDS analysis of Cr to Al ratio, the compositions of the three layers are calculated as follows: Cr0.6Al0.4N for the bonding layer, Cr0.8Al0.2N for the intermediate one and Cr0.5Al0.5N for the superficial layer. We must acknowledge that no effort was made to determine the N content. The corresponding thicknesses of the layers are ∼ 2 μm, ∼ 7 μm and ∼5 μm, respectively (see Fig. 2b). 3.2. Isothermal oxidation The squares of the mass gain as a function of oxidation time for the coated and bare alloys at 1000 °C, 1050 °C and 1100 °C are depicted in Fig. 3(a), (b) and (c), respectively. The weight gains of the coated alloy are much less than that of the bare alloy during the oxidation at the three temperatures. The straight line relationship indicates that the oxidation kinetics of the coated and bare alloy obey the parabolic rate law roughly. The parabolic rate constants (kp) have been determined and are listed in Table 3. It is apparent that for the alloy coated with the multilayer coating, the parabolic rate constants for oxidation decrease by about 2 orders of magnitude at the three temperatures. The kinetics of oxidation at 1000 °C for coated and uncoated samples during 300 h is presented in Fig. 3 (d). The mass gain of the coated sample is much less than that of the uncoated one, which approves that the coating has a good long-term resistance to oxidation. It can be clearly seen that the uncoated sample experiences a weight loss in the last ∼150 h, which indicated the breakdown of the scale. The XRD patterns of the samples after oxidation at 1000 °C, 1050 °C and 1100 °C are presented in Fig. 4. At 1000 °C, the
Fig. 7. Hot corrosion kinetics of the bare and coated alloy at (a) 900 °C, (b) 950 °C and (c) 1000 °C.
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Fig. 8. XRD patterns of the coated samples after hot corrosion in Na2SO4 salt flux at 900 °C, 950 °C and 1000 °C.
at 1000 °C, the oxide grains with disk-like shape as well as those with irregular shape appear. The grains are likely Cr2O3 and Al2O3, respectively, based on the analysis of EDS spectra shown in Fig. 5 (d) and (e). When the temperature increases to 1050 °C, the amount of irregular shape Al2O3 grains increases significantly and the disk-like Cr2O3 grains almost disappear. At 1100 °C, the size of Al2O3 grains increases slightly. The cross-section microstructure of the sample oxidized at 1050 °C is depicted in Fig. 6 together with the distribution of the elements Al, Cr, Ni, Ti and O. It can be clearly seen that an aluminum-and oxygen-rich layer, which may be Al2O3, formed on the surface of the coating and an interdiffusion between the coating and the substrate is not observed. A titanium-enriched zone is identified in the inner bonding layer of the coating, which indicates that the intermediate chromium enriched layer limits effectively the outward diffusion of Ti atoms. An oxygen rich layer forms at the inner bonding layer, which was resulted from the oxygen diffusion via the defects, such as crack, in the intermediate Cr rich layer. However, the oxygen was not observed to transport into the substrate alloy, they are trapped by the aluminum atoms in the inner bonding layer. 3.3. Hot corrosion The hot corrosion kinetics of the coated samples and the bare K38G alloy at 900 °C, 950 °C and 1000 °C are presented in Fig. 7(a), (b) and (c), respectively. When the bare alloy is exposed to oxidation at 900 °C and 950 °C, two periods can be identified from the hot corrosion kinetics curves. These two periods have been named an incubation period and a propagation period, respectively. The hot corrosion kinetics of the bare alloy exposed at 1000 °C can be divided into three periods, which are labeled as incubation, propagation and stabilization periods. However, all the coated samples exposed at 900 °C, 950 °C and 1000 °C exhibit only the incubation period during the whole exposed time of 20 hours. The kinetics curves of the coated samples indicate that the incubation period of the coated samples in the molten Na2SO4 flux is longer than 20 h at the three selected temperatures.
The XRD patterns of the coated samples after hot corrosion in Na2SO4 salt fluxing at 900 °C, 950 °C and 1000 °C are shown in Fig. 8. The phase compositions of the scales are mainly Cr2O3 with a trace of α-Al2O3. Fig. 9(a) and (b) reveal the surface morphology of the coated samples after hot corrosion at 900 °C and 1000 °C, respectively. A two-layer microstructure of the corrosion products can be distinguished on the cross-section of the samples, i.e., a loose layer on the top of the samples and a compact layer beneath it. The loose layer on the top of the sample is partly bonded with the underneath compact layer. The enlarged images of the compact bottom layer and of the loose top layer after hot corrosion tests at 1000 °C are given in Fig. 9 (c) and (d), respectively. The EDS spectra of Fig. 9 (c) and (d) are shown in Fig. 9 (e) and (f), respectively. The phase compositions of the loose layer and the compact layer were determined as α-Al2O3 and Cr2O3, respectively, based on the analysis of EDS and XRD. The cross-section microstructure as well as the distribution of the elements Al, Ti, Cr, O, Ni and S of the bare and coated alloy after exposition to hot corrosion in fused Na2SO4 at 1000 °C are shown in Fig. 10(a) and (b), respectively. One can see that the outer layer enriched in Al has been dissolved and almost completely consumed by the molten Na2SO4 salt, whereas the intermediate layer rich in Cr atoms is compact and provides a good protection against the molten salt. In comparison with the coated sample the atoms of sulfur (S) penetrate into the bare alloy more deeply, and the grain boundaries and other defects in the substrate alloy seem to act as short-circuits for transport of the sulfur atoms. 4. Discussion 4.1. Isothermal oxidation processes The outer layer of the multilayer coating has a high aluminum content, which was designed to resist the high temperature oxidation. Similar to the single-layer CrAlN coatings with a high aluminum content investigated in a previous work [22], an alumina-enriched scale can form at elevated temperatures due to the high aluminum activity (content) in the coating and the greater negative value of the Gibbs energy for the formation of alumina than that for chromia. In other words, an alumina-rich scale formation is preferential both from the kinetics as well as from the thermodynamics point of view. This is why the coating gives a good protection against oxidation at elevated temperatures. The chromium-rich intermediate layer behaves as a diffusion barrier to limit the outward diffusion of the atoms of the substrate alloy elements to the outer Al-rich layer. It is also believed that the chromium-rich interlayer can limit the loss of aluminum by diffusing into the substrate. No apparent interdiffusion is observed between the coating and the substrate, as shown in Fig. 6. Therefore, the aluminum depletion due to the interdiffusion at the interface between the coating and the substrate is minimized. As a result the high alumina content in the scale and its beneficial anti-oxidizing properties can be maintained for a long time. It is expected that the lifetime of the coated components in gas turbine engines would be enhanced.
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Fig. 9. Surface morphology of the coated samples after hot corrosion at (a) 900 °C, (b) 1000 °C. (c) and (d) are enlarged surface images of the compact bottom layer and the loose top layer after hot corrosion test at 1000 °C, respectively. (e) and (f) are the EDS spectra of (c) and (d), respectively.
4.2. Hot corrosion processes At temperatures below 1000 °C, the components can suffer another type of aggressive attack from the environment, the hot corrosion. During the incubation period, the mass gain is
relatively low and, the oxides on the surface do not break down. Once the oxide scale breaks down and its self-healing is no longer possible, the propagation period begins and a rapid deterioration of the coating takes place. Due to a lower resistance of alumina to the salt fluxing in comparison to that of
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Fig. 10. Cross-sectional microstructure of the bare (a) and coated sample (b) together with the distribution of elements Al, Ti, Cr, O, Ni and S after exposure in fused Na2SO4 at 1000 °C.
chromia, the outer Al-rich layer of the multilayer coating does not protect against the salt fluxing effectively and will dissolve in the molten salts, and precipitate as loose oxides at the interface between the salt flux and the gas environment. As a result the outer Al-rich layer is consumed rapidly and the intermediate Cr-rich layer exposes to the salt fluxing. The intermediate chromium-rich layer can readily form a continuous chromia scale and provide a rapid crack-healing route preventing the breakdown of scale. Since the coating provides protective surface oxide scales with the crack-healing properties, the incubation period can be extended greatly and the initiation of the propagation period can be increased to a great extent. This mechanism is crucial for the design of the long
lifetime components. As shown in Fig. 7, the propagation periods were not observed for all the coated samples through the exposed time in the present work, which indicates that the incubation periods for the coated samples are extended over 20 h. In comparison with the bare substrate, the incubation period of the coated alloy increases several times. For the samples exposed at 1000 °C, the incubation period of the coated sample is more than 10 times longer than that of the bare substrate. Moreover, from Fig. 7 (a), (b) and (c) it can be clearly seen that the coated samples demonstrate a slight mass loss, which was resulted from the volatilization of chromia. Erdös et al. [25] observed the mass loss during the hot corrosion of IN 939 in
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molten (Na0.9, K0.1)2SO4 at 900 °C. They assumed that the volatilization took place by the formation and subsequent decomposition of sodiumdichromate to chromia. In the present work, N2 emitting also contributed to the mass loss. From the cross-section microstructure and from the analysis of the distribution of the elements (as shown in Fig. 10(a) and (b)), the penetration of sulfur (S) is very deep for the bare substrate. In contrast, no apparent penetration of sulfur (S) was observed for the coated substrate, which indicates that the multilayer coating prevents effectively the inward diffusion of S atoms into the substrate. The results in the present paper demonstrate that the multilayer coating can respond in a pseudo-intelligent mode to the aggressive attackes of the environment, i.e., to oxidation at elevated temperatures as well as to hot corrosion at intermediate temperatures. Therefore, the multilayer CrAlN coating can be regarded as a novel type of smart coatings. 5. Conclusion A multilayer CrAlN coating of Cr0.58Al0.42N/Cr0.84Al0.16N/ Cr0.51Al0.49N was synthesized by a magnetron sputtering method on a K38G alloy. The isothermal oxidation at 1000 °C, 1050 °C and 1100 °C and the hot corrosion in a Na2SO4 salt fluxing at 900 °C, 950 °C and 1000 °C are investigated in air. Zhu [23] has investigated the corrosion behaviors of CrAlN coatings at similar temperatures. The present work demonstrated comparative resistance to oxidation but much more superior resistance to hot corrosion. The following conclusions can be drawn from the investigations performed in the present work: (1) The resistance to isothermal oxidation of the nickel-based K38G substrate alloy is enhanced significantly by the multilayer coating. The parabolic rate constants for the coated samples decrease by 2 orders of magnitude at the temperatures of 1000 °C and 1100 °C. The outer Al-rich layer contributes to the good resistance to isothermal oxidation due to formation of a continuous alumina scale on the coating's surface. (2) The intermediate Cr-rich layer can provide a diffusion barrier for the outward diffusion of the alloy elements from the substrate to the coating and the inward diffusion of Al from the coating to the substrate. (3) The resistance to hot corrosion is improved when the alloy is coated by the multilayer coating. The incubation
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period is extended from approximately 3 hours for the uncoated alloy to over 20 hours for the coated samples. The intermediate Cr-rich layer is responsible for the good resistance to the hot corrosion in molten Na2SO4 salt flux as a continuous chromia scale forms on the surface and provides a rapid route for healing the breakdown of the scale. Acknowledgements This work was supported by the National Outstanding Young Scientist Foundation of China for Y. C. Zhou under grant No. 59925208, Natural Science Foundation of China under grant No. 50232040, 50571106, 50072034, 90403027, 863 project, and High-tech Bureau of the Chinese Academy of Sciences. References [1] [2] [3] [4] [5] [6] [7] [8] [9] [10] [11] [12] [13] [14] [15] [16] [17] [18] [19] [20] [21] [22] [23] [24] [25]
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