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High temperature behavior of a diffusion barrier coating evolved from ZrO2 precursor layer
Surface & Coatings Technology 357 (2019) 384–392
Contents lists available at ScienceDirect
Surface & Coatings Technology journal homepage: www.elsevier.com/locate/surfcoat
High temperature behavior of a diffusion barrier coating evolved from ZrO2 precursor layer
T
Chenxi Yanga, Zhengxian Lib, Lintao Liub, Fan Yea, Sujun Wua,
⁎
a b
School of Materials Science and Engineering, Beihang University, Beijing 100191, China Department of Corrosion and Protection, Northwest Institute for Non-ferrous Metal Research, Xi'an 710016, China
ARTICLE INFO
ABSTRACT
Keywords: Ni-based superalloy ZrO2 Diffusion barrier Interdiffusion
The zirconia (ZrO2) as a precursor layer for the diffusion barrier coating (DBC) was deposited by electron beam physical vapor deposition (EB-PVD) between the René N5 superalloy substrate and NiCrAl coating. The high temperature exposure for the N5/ZrO2/NiCrAl system was carried out at 1000 °C for 5 h, 100 h and 200 h in atmosphere to investigate the structure evolution of the ZrO2 precursor layer and the elements interdiffusion of the system compared with the N5/NiCrAl system. The results showed that a sandwich structural DBC (α-Al2O3/ rich-Zr/α-Al2O3) was formed through redox reaction between zirconia and aluminum under high temperature, leading to the transformation of N5/ZrO2/NiCrAl system into the N5/DBC/NiCrAl system. It was found that the sandwiched DBC could effectively suppress the interdiffusion of elements compared to the N5/NiCrAl system.
1. Introduction The single-crystalline Ni-based superalloys as high-temperature structure materials have been widely applied in aerospace industry due to their excellent mechanical performance, especially the creep strength [1–3]. The outstanding performance is attributed to the increase of refractory metal elements, such as tantalum, rhenium, tungsten, and molybdenum [4]. However, the oxidation and corrosion resistance of the superalloys would deteriorate due to the reduction of Al and Cr relative contents. MCrAlX (M: Ni, Co or Ni + Co, X: Y, Si, Hf) coatings have been used to protect Ni-based superalloys in high-temperature oxidation and corrosion environments. However, some problems would arise due to the elements interdiffusion between the MCrAlX coatings and the Nibased superalloy substrate during high temperature service. The continuous growth of protective oxide film (Al2O3) on the MCrAlX coating surface would be hindered by the inward diffusion of Al and Cr from coating to substrate and the outward diffusion of the refractory elements. The Kirkendall voids could be produced through elements diffusion at the interface between the substrate and the coating, which could weaken the bonding force of substrate and MCrAlX coating [5]. The second reaction zone (SRZ) and topological close packed (TCP) phases in the substrate would be formed, resulting in the degradation of the creep strength [6]. In order to prevent the elements interdiffusion, many studies have
⁎
been reported, such as equilibrium coatings (EQ coatings) and diffusion barrier coatings. The EQ coatings were developed to be in thermodynamic equilibrium with the underlying Ni-based substrate, which could minimize the elements interdiffusion and suppress the formation of SRZ due to the equal chemical potentials of the alloying elements between the coating and substrate [7–9]. Diffusion barrier coatings are mainly divided into two types, metals and ceramics. The metallic diffusion barriers, i.e. Ru [10], Re-Cr-Ni [11], NieW [12], have good bonding strength between Ni-base superalloy substrate and MCrAlX coating, but have poor blocking ability of elements interdiffusion. The ceramic diffusion barriers, i.e. AlN [13] and α-Al2O3 [14], have a good ability to prevent interdiffusion but have the weak bonding strength due to their much lower thermal expansion coefficients than that of the substrate and MCrAlX coatings [15]. Therefore, the active diffusion barriers have been proposed by researchers [16–19]. As a source of oxygen, a precursor oxide layer with the lower thermodynamic stability than α-Al2O3, such as Cr2O3 [16], Cr-O-N [17,18] and yttria partially stabilized zirconia [19], was applied to the interface between the substrate and MCrAlX coating to form a sandwich structural diffusion barrier coating (α-Al2O3/rich-Metal/α-Al2O3) through redox reaction. This sandwich structure can effectively prevent the elements interdiffusion, and also possesses high reaction bonding strength. In this paper, the pure zirconia (ZrO2) as a precursor layer of the active diffusion barriers was deposited between the N5 superalloy substrate and NiCrAl coating by electron beam physical vapor
Corresponding author. E-mail address: [email protected] (S. Wu).
https://doi.org/10.1016/j.surfcoat.2018.10.022 Received 26 June 2018; Received in revised form 8 October 2018; Accepted 9 October 2018 Available online 09 October 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.
Surface & Coatings Technology 357 (2019) 384–392
C. Yang et al.
were analyzed compared to the system without DBC. 2. Experimental procedures The single-crystalline Ni-based superalloy René N5 (nominal chemical compositions: 7 wt% Cr, 7.5 wt% Co, 5.0 wt% W, 6.2 wt% Al, 6.5 wt% Ta, 1.5 wt% Mo, 3.0 wt% Re, 0.15 wt% Hf, Ni: Bal.) was used as the substrate in this study. The N5 superalloy was cut into cylindershape with 10 mm diameter and 3 mm thickness. The surface of the specimens was cleaned by ultrasonic within ethanol and acetone after polishing. The N5/ZrO2/NiCrAl system was manufactured by electron beam physical vapor deposition (EB-PVD). The thin ZrO2 precursor layer (purity: 99.9%) was deposited on the substrate at 800 °C, and then the Ni-20Cr-10Al coating was deposited on the ZrO2 layer. The electron beam current for evaporating the ZrO2 and NiCrAl targets was set at ~1.2A and ~1A, respectively. During the deposition, the vacuum of the evaporation chamber was kept at ~10−3 Pa. The deposited thickness of the ZrO2 precursor layer was 2 μm and that of the deposited NiCrAl coating was 15 μm. The N5/ZrO2 system was produced to explore whether the ZrO2 was dissolved during the deposition. The N5/NiCrAl system was also prepared for comparison. After deposition, the high temperature exposure for N5/ZrO2/ NiCrAl system was carried out at the temperature of 1000 °C for 5 h, 100 h, and 200 h in air to observe the structure evolution of the ZrO2 precursor layer. The interdiffusion and oxidation behaviors of the system were investigated by comparing with the N5/NiCrAl system at 1000 °C for 200 h. The roughness of the substrate was characterized by Atomic Force Microscope (AFM, DIMENSION ICON, Bruker, USA). The surface and cross-section morphologies of the coatings were observed by scanning electron microscopy (SEM, JSM-6460, JEOL Ltd., Japan) equipped with the energy-dispersive spectrometer (EDS, X-Sight, Oxford instruments Co, Oxford, UK). The chemical bonding states of elements were analyzed by X-ray photoelectron spectroscopy (XPS, Thermo Scientific ESCALAB250Xi, USA). The X-ray diffraction (XRD, D/MAX-2500, Rigaku, Japan) tests were performed to identify the phase constitution.
Fig. 1. AFM images of N5 superalloy substrates after polishing processing.
3. Results and discussion 3.1. Microstructural characterization of the as-deposited samples Before depositing the ZrO2 precursor layer, the surface roughness of the substrate was characterized by AFM. As shown in Fig. 1, the scanning area was 50 μm × 50 μm in every different regions of the substrate surface to calculate the mean value of Ra, and finally the obtained Ra was 13.1 nm. Fig. 2 shows the surface morphology of the precursor layer in asdeposited N5/ZrO2 system and corresponding XRD pattern. As shown in Fig. 2a, the rough surface is composed of fine grains. The XRD pattern (Fig. 2b) display that the component of the precursor layer was c-ZrO2, which indicated that ZrO2 was not decomposed during deposition. Fig. 3 represents the surface image and XRD pattern of the NiCrAl coating. As shown in Fig. 3a, the surface of the NiCrAl coating presents a relative smooth surface. The XRD pattern (Fig. 3b) indicates that γ´Ni3Al, β-NiAl, and α-Cr were precipitated in as-deposited NiCrAl coating. According to the ICDD database, β-NiAl and α-Cr are bcc lattice structures with close lattice parameters (a = 0.2880 nm in β-NiAl and a = 0.2879 nm in α-Cr). Therefore, it is difficult to distinguish the two phases in XRD pattern. It was reported that, at ambient
Fig. 2. Surface (a) and XRD pattern (b) of the as-deposited ZrO2 coating.
deposition (EB-PVD). The structure evolution of the ZrO2 precursor layer was investigated at 1000 °C with different exposure time. The interdiffusion and oxidation behaviors of the N5/DBC/NiCrAl system
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Fig. 3. Surface (a), XRD pattern (b) and EDS spectrum (c) of the as-deposited NiCrAl coating.
temperature, the solid solubility limit of Cr in γ`-Ni3Al and β-NiAl is 4 at.% and 20 at.% respectively [20]. In this study, however, the content of Cr was measured as 27 at.% in the coating surface (Fig. 3c), indicating the existence of α-Cr although its diffraction peak was completely overlapped with that of the β-NiAl. The cross-sectional morphologies of the N5/NiCrAl system are shown in Fig. 4a. The interface was continuous and smooth between N5 substrate and NiCrAl coating. The microscopic interdiffusion zone (IDZ) was formed due to the Ni, Cr, Al elements diffusion from NiCrAl coating to N5 substrate (see EDS results in Fig. 4b). Fig. 5a shows the cross-sectional morphologies of the N5/ZrO2/ NiCrAl system. It can be found that a thin black layer was generated at the interface of the ZrO2/NiCrAl through reaction between ZrO2 and the NiCrAl coating during the depositing process. The EDS results (Fig. 5b) indicated that the content of the Al is much higher than other elements in the black layer. According to the thermodynamics data [21], the standard Gibbs free energy varied with the temperature for NiO, Cr2O3, Al2O3, as shown in Fig. 6, and the order of the oxides stability can be described as Al2O3 > Cr2O3 > NiO. The results demonstrated that the affinity of Al with O is higher than that of Ni and Cr. XPS was used to further investigate the chemical bonding states of Al and Zr elements in the DBC. As shown in the XPS Al2p spectra (Fig. 7a), a peak exists at a binding energy of 74.3 eV, which is consistent with the binding energy of Al2p in α-Al2O3. Therefore, the black layer can be considered as Al2O3. The Zr3d peaks were fitted into 4 peaks in Fig. 7b, the peaks at binding energies of 181.1 eV and 184.4 eV were consistent with the binding energy of Zr3d in ZrO2, and the peaks
at binding energies of 178.6 eV and 179.6 eV indicate the formation of free Zr. But the content of Zr was relatively low due to the slow reaction at as-deposited state. The possible reaction between Al and ZrO2 is as follows:
4Al + 3ZrO2
2Al2 O3 + 3[Zr]
(1)
where [Zr] stands for the free Zr, and the standard Gibbs free energy of the reaction is expressed by:
G0 =
176,603 + 54.63T
(2)
0
when ΔG = 0, T is 2960 °C. The variation of the standard Gibbs free energy with the temperature for the reaction is also shown in Fig. 6. The heat exposure temperature in this experiment is lower than 2960 °C, thus ΔG0 < 0, which indicates the reaction can take place spontaneously to form alumina. 3.2. Evolution of the ZrO2 precursor layer In order to study the structure evolution of the ZrO2 precursor layer, the high temperature exposure tests were conducted at 1000 °C for different exposure time (5 h, 100 h, and 200 h). After 5 h exposure, the thin black layer of α-Al2O3 appeared at the interface of the N5/ZrO2 (Fig. 8a) and the black layer at ZrO2/NiCrAl became thicker, resulting in a sandwich structural DBC (α-Al2O3/rich-Zr/α-Al2O3). As shown in Figs. 6a–8c, the α-Al2O3 layers became thicker and the rich-Zr layer became narrower with the increase of the high temperature exposure time. The variation trend of the thickness for each layer in the sandwich
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Fig. 5. Cross-sectional morphologies (a) of the as-deposited coatings: N5/ZrO2/ NiCrAl system and EDS line-scan result (b).
Fig. 4. Cross-sectional morphologies (a) of the as-deposited N5/NiCrAl system and EDS line-scan result (b).
along the thickness of the N5/NiCrAl coating system, which was exposed at 1000 °C for 200 h in the atmosphere. As shown in Fig. 9a, the IDZ is thicker than that of the as-deposited specimen (Fig. 4a). The
structure was quantified in Fig. 8d. The average thicknesses of the αAl2O3 layers close to the substrate and the NiCrAl coating are 0.46 μm and 0.72 μm for 5 h,1.09 μm and 1.44 μm for 100 h,1.22 μm and 2.11 μm for 200 h, respectively. The thicknesses of the rich-Zr layer are 1.61 μm (5 h), 1.39 μm (100h), and 1.03 μm (200h). The results also indicated that the α-Al2O3 layers can still grow through the diffusion of the Al and/or O atoms with the increase of exposure time. Additionally, the α-Al2O3 layer close to the NiCrAl is thicker than that close to the substrate during high temperature exposure, which may be attributed to the following factors: (1) Al content in the substrate (6.2 wt%) is lower than that in NiCrAl coating (10 wt%); (2) The activation and migration energies of Al atoms in β-NiAl phase are lower than those in γ´-Ni3Al phase [22]; (3) the refractory elements in substrate (i.e., W and Mo) can improve the bonding force between Ni and Al atoms, leading to the decrease of diffusion rate; (4) grain boundaries in NiCrAl coating can promote elements diffusion, but there are no grain boundaries in the single-crystalline substrate. 3.3. Diffusion resistance of the diffusion barrier
Fig. 6. Standard Gibbs free energy for Al2O3, Cr2O3, NiO and [Zr] as function of temperature.
Fig. 9 shows the cross-sectional morphology and element profiles
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along the thickness of the N5/DBC/NiCrAl system that exposed at 1000 °C for 200 h in the atmosphere. It can be found that the α-Cr phase has not formed in the NiCrAl coating (Fig. 10a). The IDZ, SRZ, and voids were not formed compared with the N5/NiCrAl system, and a good adherence of the interface was observed. The elements profiles along the thickness of the N5/DBC/NiCrAl system were measured as shown in Fig. 10b–j. The selected areas for EDS are labeled with number (points 1 to 7 in Fig. 10a). The O and Zr were only detected in the sandwiched diffusion barrier. The diffusion of Ni and Cr cannot be observed from Fig. 10d and e. A small amount of Co was found in the α-Al2O3 layer adjacent to substrate side, which may be attribute that the alumina diffusion barrier was not dense in the initial formation stage, so that the Co was diffused to the α-Al2O3 layer. From Fig. 10h–j, it can be seen that the diffusion of the refractory elements W, Ta, and Re was blocked by the DBC. The aluminum as an importantly diffused element, its contents from point 1 to 7 were 3.12, 2.04, 31.03, 8.85, 50.17, 0 and 0 wt%, respectively. It can be noted that the depletion of Al occurred in the NiCrAl coating close to the DBC, which would suppress the growth of the α-Al2O3 in the diffusion barrier coating. In the N5/NiCrAl system, the elements interdiffusion were caused due to the inward diffusion of the alloying elements from coating to substrate and the outward diffusion of the refractory elements from substrate to NiCrAl coating. The contents of the refractory elements were relatively low in the substrate, which resulted in the low concentration gradient along the diffusion direction. As is well known, the diffusion coefficient of the refractory elements is normally lower than that of the relatively low melting metal elements. Therefore, the contents and diffusion coefficients of the alloying elements in NiCrAl coating are higher than that of the refractory elements in the substrate. These results suggest that the diffusion fluxes of atoms from NiCrAl coating to substrate were higher than that from substrate to coating, leading to the formations of the IDZ, SRZ and Kirkendall porosities. Additionally, the diffusion of Al was the main factor of affecting the performance of the coating and substrate during the elements interdiffusion. According to the first principal of Fick's law:
Fig. 7. XPS spectra of the Al element (a) and Zr element (b) in the interface reaction zone between ZrO2/NiCrAl coating under as-deposited state.
formation of the IDZ was resulted from the interdiffusion of Al, Cr, and Ni at the interface of substrate and NiCrAl coating, which can promote the precipitation of refractory elements (e.g. Re, W, and Ta) to form SRZ zone and needle-like TCP phase in substrate. In addition, some dark phases were precipitated in the NiCrAl coating, and some Kirkendall voids were observed at the IDZ/NiCrAl interface. Elemental distribution along thickness direction for the N5/NiCrAl system was measured by EDS (Fig. 9b–h). The trace to measure the elemental profile is a row of points (numbered 1–6) shown in Fig. 9a. Variation of the Ni, Cr, Al, and Co elements in the IDZ can be ascribed to elemental interdiffusion. The high content of Cr (87.4 in wt%) in the dark phase (point 4) suggests that the phase should be α-Cr. A certain amount of W (~2.7%) was detected in the α-Cr phase (Fig. 9e), implying that W might have played some role in the formation of α-Cr. The existence of Co (Fig. 9f) and W elements in the NiCrAl coating indicated that these two elements diffused from the substrate across the interface into the NiCrAl coating layer. Fig. 10 shows the cross-sectional morphology and element profiles
J=
D
d dx
(3)
where J is the diffusion flux, D is the diffusion coefficient, ρ is the mass concentration of the diffused substance, and dρ/dx is the concentration gradient along the thickness direction. The diffusion flux of Al is proportional to the diffusion coefficient and the mass concentration, but inversely proportional to the diffusion distance that can be regarded as the thickness of the diffusion barrier. For the N5/ZrO2/NiCrAl system, the initial single ZrO2 precursor layer will transform into a sandwich structure when Al reacts with ZrO2 continuously to form Al2O3. That led to the increase of diffusion barrier thickness, i.e., the diffusion distance. Simultaneously, the alloying elements have the lower diffusion rate in α-Al2O3 than that in the DBC and NiCrAl coating [17]. Therefore, the interdiffusion of the metal elements would become much more difficult, and the formation of IDZ, SRZ and voids were suppressed due to the existence of the DBC.
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Fig. 8. Cross-sectional morphology of diffusion barrier exposed at 1000 °C for 5 h (a), 100 h (b), 200 h (c) and the variation of average thickness with the oxidation time (d).
3.4. Surface morphologies of the heat-treated samples
interdiffusion of the two systems were studied. Conclusions can be drawn as follows:
The surface morphologies of NiCrAl coatings in the N5/NiCrAl and N5/DBC/NiCrAl systems were shown in Fig. 11. The spallation of the scales can be observed in the two systems. The XRD patterns of the NiCrAl surfaces in the two systems after 200 h exposure at 1000 °C in atmosphere were shown in Fig. 12. As can be seen, the surfaces of the two coatings were composed of γ-Ni and α-Al2O3. The results revealed that the original β-NiAl and the γ´-Ni3Al phases in Fig. 3 have transformed into the γ-Ni phase because of the formation of the α-Al2O3. The generation of α-Al2O3 scale in surfaces could improve the oxidation resistance of the systems.
(1) The interdiffusion zone (IDZ), secondary reaction zone (SRZ) and Kirkendall voids were formed in the N5/NiCrAl system, due to the interdiffusion of elements after thermal exposure at 1000 °C for 200 h. (2) For the N5/ZrO2/NiCrAl system, the ZrO2 precursor layer transformed into a sandwich structural DBC (α-Al2O3/rich-Zr/α-Al2O3) after thermal exposure at 1000 °C, and the average thickness of each α-Al2O3 layer increased with the increase of thermal exposure time. (3) The DBC could suppress the interdiffusion of the alloying elements compared to the system without the diffusion barrier coating.
4. Conclusions The effect of thermal exposure on the structure evolution of the N5/ NiCrAl and N5/ZrO2/NiCrAl systems was analyzed. The elements
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Fig. 9. Cross-sectional morphology (a) and elements profiles (b–h) along the thickness of the N5/NiCrAl system after 200 h exposure at 1000 °C in atmosphere.
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Fig. 10. Cross-sectional morphologies (a) and element profiles (b-j) along the thickness of the N5/DBC/NiCrAl system after 200 h exposure at 1000 °C in atmosphere.
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Acknowledgments This work was supported by National Key Basic Research Program of China (973 Program, No. 2012CB625102). References [1] Q. Lu, Y.L. Pei, S.S. Li, et al., Role of volatilization of molybdenum oxides during the cyclic oxidation of high-Mo containing Ni-based single crystal superalloys, Corros. Sci. 129 (2017) 192–204. [2] N.D. Souza, J. Kelleher, C.L. Qiu, et al., The role of stress relaxation and creep during high temperature deformation in Ni-base single crystal superalloys - implications to strain build-up during directional solidification, Acta Mater. 106 (2016) 322–332. [3] R. Farangis, Z.M. Li, S. Zaefferer, et al., On the origin of creep dislocations in a Nibase, single-crystal superalloy: an ECCI, EBSD, and dislocation dynamics-based study, Acta Mater. 109 (2016) 151–161. [4] M.J. Pomeroy, Coatings for gas turbine materials and long term stability issues, Mater. Des. 26 (2005) 223–231. [5] M. Elsaß, M. Frommherz, A. Scholz, et al., Interdiffusion in MCrAlY coated nickelbase superalloys, Surf. Coat. Technol. 307 (2016) 565–573. [6] H.R. Yao, Z.B. Bao, M.L. Shen, et al., A magnetron sputtered microcrystalline β-NiAl coating for SC superalloys. Part II. Effects of a NiCrO diffusion barrier on oxidation behavior at 1100 °C, Appl. Surf. Sci. 407 (2017) 485–494. [7] A. Sato, H. Harada, K. Kawagishi, Development of a new bond coat “EQ coating” system, Metall. Mater. Trans. A 37A (2006) 789–790. [8] K. Kawagishi, A. Sato, H. Harada, A concept for the EQ coating system for nickelbased superalloys, JOM-USL 60 (2008) 31–35. [9] K. Kawagishi, H. Harada, Akihiro Sato, et al., EQ coating: a new concept for SRZfree coating systems, Superalloys (2008) 761–768. [10] Z. Bai, D. Li, H. Peng, et al., Suppressing the formation of SRZ in a Ni-based single crystal superalloy by RuNiAl diffusion barrier, Prog. Nat. Sci. 22 (2012) 146–152. [11] T. Narita, F. Lang, K. Thosin, et al., Oxidation behavior of nickel-base single-crystal superalloy with rhenium-base diffusion barrier coating system at 1,423 K in air, Oxid. Met. 68 (2007) 343–363. [12] E. Cavaletti, S. Naveos, S. Mercier, et al., Ni–W diffusion barrier: its influence on the oxidation behaviour of a β-(Ni, Pt) Al coated fourth generation nickel-base superalloy, Surf. Coat. Technol. 204 (2009) 761–765. [13] L.J. Zhu, S.L. Zhu, F.H. Wang, Preparation and oxidation behavior of nanocrystalline Ni+CrAlYSiN composite coating with AlN diffusion barrier on Ni-based superalloy K417, Corros. Sci. 60 (2012) 265–274. [14] S.H. Hosseini, S. Mirdamadi, S. Rastegari, Investigating efficiency of α-Al2O3 diffusion barrier layer in oxidation of EB-PVD NiCrAlY coatings, Surf. Eng. 31 (2015) 146–155. [15] P. Ren, S.L. Zhu, F.H. Wang, Microstructural stability of AlN diffusion barrier for nanocompositeNi + CrAlYSiHfN coating on single crystal superalloy at high temperatures, Appl. Surf. Sci. 359 (2015) 420–425. [16] Y.X. Xu, X.T. Luo, C.X. Li, Effect of annealing conditions on in situ formation of Cr2O3 diffusion barrier, Surf. Eng. 33 (2017) 210–216. [17] Q.M. Wang, Y.N. Wu, M.H. Guo, et al., Ion-plated Al-O-N and Cr-O-N films on Nibase superalloys as diffusion barriers, Surf. Coat. Technol. 197 (2005) 68–76. [18] W.Z. Li, Q.M. Wang, Z.B. Bao, et al., Microstructural evolution of the NiCrAlY/CrON duplex coating system and its influence on mechanical properties, Mater. Sci. Eng. A 498 (2008) 487–494. [19] C.A. Guo, W. Wang, Y.X. Cheng, et al., Yttria partially stabilised zirconia as diffusion barrier between NiCrAlY and Ni-base single crystal René N5 superalloy, Corros. Sci. 94 (2015) 122–128. [20] S. Merchant, M.R. Notis, A review – constitution of the Al–Cr–Ni system, Mater. Sci. Eng. 66 (1984) 47–60. [21] I. Barin, O. Knacke, O. Kubaschewski, Thermochemical Properties of Inorganic Substances, Springer-Verlag, Berlin, 1991. [22] S. Yu, T. Yu C Wang, et al., Self-diffusion in the intermetallic compounds NiAl and Ni3Al-An embedded atom method study, Physica B 396 (2007) 138–144.
Fig. 11. Surface morphologies of the NiCrAl after 200 h exposure at 1000 °C in atmosphere: N5/NiCrAl system (a) and N5/DBC/NiCrAl system (b).
Fig. 12. XRD patterns of the NiCrAl surface after 200 h exposure at 1000 °C in air.
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High temperature behavior of a diffusion barrier coating evolved from ZrO2 precursor layer
T
Chenxi Yanga, Zhengxian Lib, Lintao Liub, Fan Yea, Sujun Wua,
⁎
a b
School of Materials Science and Engineering, Beihang University, Beijing 100191, China Department of Corrosion and Protection, Northwest Institute for Non-ferrous Metal Research, Xi'an 710016, China
ARTICLE INFO
ABSTRACT
Keywords: Ni-based superalloy ZrO2 Diffusion barrier Interdiffusion
The zirconia (ZrO2) as a precursor layer for the diffusion barrier coating (DBC) was deposited by electron beam physical vapor deposition (EB-PVD) between the René N5 superalloy substrate and NiCrAl coating. The high temperature exposure for the N5/ZrO2/NiCrAl system was carried out at 1000 °C for 5 h, 100 h and 200 h in atmosphere to investigate the structure evolution of the ZrO2 precursor layer and the elements interdiffusion of the system compared with the N5/NiCrAl system. The results showed that a sandwich structural DBC (α-Al2O3/ rich-Zr/α-Al2O3) was formed through redox reaction between zirconia and aluminum under high temperature, leading to the transformation of N5/ZrO2/NiCrAl system into the N5/DBC/NiCrAl system. It was found that the sandwiched DBC could effectively suppress the interdiffusion of elements compared to the N5/NiCrAl system.
1. Introduction The single-crystalline Ni-based superalloys as high-temperature structure materials have been widely applied in aerospace industry due to their excellent mechanical performance, especially the creep strength [1–3]. The outstanding performance is attributed to the increase of refractory metal elements, such as tantalum, rhenium, tungsten, and molybdenum [4]. However, the oxidation and corrosion resistance of the superalloys would deteriorate due to the reduction of Al and Cr relative contents. MCrAlX (M: Ni, Co or Ni + Co, X: Y, Si, Hf) coatings have been used to protect Ni-based superalloys in high-temperature oxidation and corrosion environments. However, some problems would arise due to the elements interdiffusion between the MCrAlX coatings and the Nibased superalloy substrate during high temperature service. The continuous growth of protective oxide film (Al2O3) on the MCrAlX coating surface would be hindered by the inward diffusion of Al and Cr from coating to substrate and the outward diffusion of the refractory elements. The Kirkendall voids could be produced through elements diffusion at the interface between the substrate and the coating, which could weaken the bonding force of substrate and MCrAlX coating [5]. The second reaction zone (SRZ) and topological close packed (TCP) phases in the substrate would be formed, resulting in the degradation of the creep strength [6]. In order to prevent the elements interdiffusion, many studies have
⁎
been reported, such as equilibrium coatings (EQ coatings) and diffusion barrier coatings. The EQ coatings were developed to be in thermodynamic equilibrium with the underlying Ni-based substrate, which could minimize the elements interdiffusion and suppress the formation of SRZ due to the equal chemical potentials of the alloying elements between the coating and substrate [7–9]. Diffusion barrier coatings are mainly divided into two types, metals and ceramics. The metallic diffusion barriers, i.e. Ru [10], Re-Cr-Ni [11], NieW [12], have good bonding strength between Ni-base superalloy substrate and MCrAlX coating, but have poor blocking ability of elements interdiffusion. The ceramic diffusion barriers, i.e. AlN [13] and α-Al2O3 [14], have a good ability to prevent interdiffusion but have the weak bonding strength due to their much lower thermal expansion coefficients than that of the substrate and MCrAlX coatings [15]. Therefore, the active diffusion barriers have been proposed by researchers [16–19]. As a source of oxygen, a precursor oxide layer with the lower thermodynamic stability than α-Al2O3, such as Cr2O3 [16], Cr-O-N [17,18] and yttria partially stabilized zirconia [19], was applied to the interface between the substrate and MCrAlX coating to form a sandwich structural diffusion barrier coating (α-Al2O3/rich-Metal/α-Al2O3) through redox reaction. This sandwich structure can effectively prevent the elements interdiffusion, and also possesses high reaction bonding strength. In this paper, the pure zirconia (ZrO2) as a precursor layer of the active diffusion barriers was deposited between the N5 superalloy substrate and NiCrAl coating by electron beam physical vapor
Corresponding author. E-mail address: [email protected] (S. Wu).
https://doi.org/10.1016/j.surfcoat.2018.10.022 Received 26 June 2018; Received in revised form 8 October 2018; Accepted 9 October 2018 Available online 09 October 2018 0257-8972/ © 2018 Elsevier B.V. All rights reserved.
Surface & Coatings Technology 357 (2019) 384–392
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were analyzed compared to the system without DBC. 2. Experimental procedures The single-crystalline Ni-based superalloy René N5 (nominal chemical compositions: 7 wt% Cr, 7.5 wt% Co, 5.0 wt% W, 6.2 wt% Al, 6.5 wt% Ta, 1.5 wt% Mo, 3.0 wt% Re, 0.15 wt% Hf, Ni: Bal.) was used as the substrate in this study. The N5 superalloy was cut into cylindershape with 10 mm diameter and 3 mm thickness. The surface of the specimens was cleaned by ultrasonic within ethanol and acetone after polishing. The N5/ZrO2/NiCrAl system was manufactured by electron beam physical vapor deposition (EB-PVD). The thin ZrO2 precursor layer (purity: 99.9%) was deposited on the substrate at 800 °C, and then the Ni-20Cr-10Al coating was deposited on the ZrO2 layer. The electron beam current for evaporating the ZrO2 and NiCrAl targets was set at ~1.2A and ~1A, respectively. During the deposition, the vacuum of the evaporation chamber was kept at ~10−3 Pa. The deposited thickness of the ZrO2 precursor layer was 2 μm and that of the deposited NiCrAl coating was 15 μm. The N5/ZrO2 system was produced to explore whether the ZrO2 was dissolved during the deposition. The N5/NiCrAl system was also prepared for comparison. After deposition, the high temperature exposure for N5/ZrO2/ NiCrAl system was carried out at the temperature of 1000 °C for 5 h, 100 h, and 200 h in air to observe the structure evolution of the ZrO2 precursor layer. The interdiffusion and oxidation behaviors of the system were investigated by comparing with the N5/NiCrAl system at 1000 °C for 200 h. The roughness of the substrate was characterized by Atomic Force Microscope (AFM, DIMENSION ICON, Bruker, USA). The surface and cross-section morphologies of the coatings were observed by scanning electron microscopy (SEM, JSM-6460, JEOL Ltd., Japan) equipped with the energy-dispersive spectrometer (EDS, X-Sight, Oxford instruments Co, Oxford, UK). The chemical bonding states of elements were analyzed by X-ray photoelectron spectroscopy (XPS, Thermo Scientific ESCALAB250Xi, USA). The X-ray diffraction (XRD, D/MAX-2500, Rigaku, Japan) tests were performed to identify the phase constitution.
Fig. 1. AFM images of N5 superalloy substrates after polishing processing.
3. Results and discussion 3.1. Microstructural characterization of the as-deposited samples Before depositing the ZrO2 precursor layer, the surface roughness of the substrate was characterized by AFM. As shown in Fig. 1, the scanning area was 50 μm × 50 μm in every different regions of the substrate surface to calculate the mean value of Ra, and finally the obtained Ra was 13.1 nm. Fig. 2 shows the surface morphology of the precursor layer in asdeposited N5/ZrO2 system and corresponding XRD pattern. As shown in Fig. 2a, the rough surface is composed of fine grains. The XRD pattern (Fig. 2b) display that the component of the precursor layer was c-ZrO2, which indicated that ZrO2 was not decomposed during deposition. Fig. 3 represents the surface image and XRD pattern of the NiCrAl coating. As shown in Fig. 3a, the surface of the NiCrAl coating presents a relative smooth surface. The XRD pattern (Fig. 3b) indicates that γ´Ni3Al, β-NiAl, and α-Cr were precipitated in as-deposited NiCrAl coating. According to the ICDD database, β-NiAl and α-Cr are bcc lattice structures with close lattice parameters (a = 0.2880 nm in β-NiAl and a = 0.2879 nm in α-Cr). Therefore, it is difficult to distinguish the two phases in XRD pattern. It was reported that, at ambient
Fig. 2. Surface (a) and XRD pattern (b) of the as-deposited ZrO2 coating.
deposition (EB-PVD). The structure evolution of the ZrO2 precursor layer was investigated at 1000 °C with different exposure time. The interdiffusion and oxidation behaviors of the N5/DBC/NiCrAl system
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Fig. 3. Surface (a), XRD pattern (b) and EDS spectrum (c) of the as-deposited NiCrAl coating.
temperature, the solid solubility limit of Cr in γ`-Ni3Al and β-NiAl is 4 at.% and 20 at.% respectively [20]. In this study, however, the content of Cr was measured as 27 at.% in the coating surface (Fig. 3c), indicating the existence of α-Cr although its diffraction peak was completely overlapped with that of the β-NiAl. The cross-sectional morphologies of the N5/NiCrAl system are shown in Fig. 4a. The interface was continuous and smooth between N5 substrate and NiCrAl coating. The microscopic interdiffusion zone (IDZ) was formed due to the Ni, Cr, Al elements diffusion from NiCrAl coating to N5 substrate (see EDS results in Fig. 4b). Fig. 5a shows the cross-sectional morphologies of the N5/ZrO2/ NiCrAl system. It can be found that a thin black layer was generated at the interface of the ZrO2/NiCrAl through reaction between ZrO2 and the NiCrAl coating during the depositing process. The EDS results (Fig. 5b) indicated that the content of the Al is much higher than other elements in the black layer. According to the thermodynamics data [21], the standard Gibbs free energy varied with the temperature for NiO, Cr2O3, Al2O3, as shown in Fig. 6, and the order of the oxides stability can be described as Al2O3 > Cr2O3 > NiO. The results demonstrated that the affinity of Al with O is higher than that of Ni and Cr. XPS was used to further investigate the chemical bonding states of Al and Zr elements in the DBC. As shown in the XPS Al2p spectra (Fig. 7a), a peak exists at a binding energy of 74.3 eV, which is consistent with the binding energy of Al2p in α-Al2O3. Therefore, the black layer can be considered as Al2O3. The Zr3d peaks were fitted into 4 peaks in Fig. 7b, the peaks at binding energies of 181.1 eV and 184.4 eV were consistent with the binding energy of Zr3d in ZrO2, and the peaks
at binding energies of 178.6 eV and 179.6 eV indicate the formation of free Zr. But the content of Zr was relatively low due to the slow reaction at as-deposited state. The possible reaction between Al and ZrO2 is as follows:
4Al + 3ZrO2
2Al2 O3 + 3[Zr]
(1)
where [Zr] stands for the free Zr, and the standard Gibbs free energy of the reaction is expressed by:
G0 =
176,603 + 54.63T
(2)
0
when ΔG = 0, T is 2960 °C. The variation of the standard Gibbs free energy with the temperature for the reaction is also shown in Fig. 6. The heat exposure temperature in this experiment is lower than 2960 °C, thus ΔG0 < 0, which indicates the reaction can take place spontaneously to form alumina. 3.2. Evolution of the ZrO2 precursor layer In order to study the structure evolution of the ZrO2 precursor layer, the high temperature exposure tests were conducted at 1000 °C for different exposure time (5 h, 100 h, and 200 h). After 5 h exposure, the thin black layer of α-Al2O3 appeared at the interface of the N5/ZrO2 (Fig. 8a) and the black layer at ZrO2/NiCrAl became thicker, resulting in a sandwich structural DBC (α-Al2O3/rich-Zr/α-Al2O3). As shown in Figs. 6a–8c, the α-Al2O3 layers became thicker and the rich-Zr layer became narrower with the increase of the high temperature exposure time. The variation trend of the thickness for each layer in the sandwich
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Fig. 5. Cross-sectional morphologies (a) of the as-deposited coatings: N5/ZrO2/ NiCrAl system and EDS line-scan result (b).
Fig. 4. Cross-sectional morphologies (a) of the as-deposited N5/NiCrAl system and EDS line-scan result (b).
along the thickness of the N5/NiCrAl coating system, which was exposed at 1000 °C for 200 h in the atmosphere. As shown in Fig. 9a, the IDZ is thicker than that of the as-deposited specimen (Fig. 4a). The
structure was quantified in Fig. 8d. The average thicknesses of the αAl2O3 layers close to the substrate and the NiCrAl coating are 0.46 μm and 0.72 μm for 5 h,1.09 μm and 1.44 μm for 100 h,1.22 μm and 2.11 μm for 200 h, respectively. The thicknesses of the rich-Zr layer are 1.61 μm (5 h), 1.39 μm (100h), and 1.03 μm (200h). The results also indicated that the α-Al2O3 layers can still grow through the diffusion of the Al and/or O atoms with the increase of exposure time. Additionally, the α-Al2O3 layer close to the NiCrAl is thicker than that close to the substrate during high temperature exposure, which may be attributed to the following factors: (1) Al content in the substrate (6.2 wt%) is lower than that in NiCrAl coating (10 wt%); (2) The activation and migration energies of Al atoms in β-NiAl phase are lower than those in γ´-Ni3Al phase [22]; (3) the refractory elements in substrate (i.e., W and Mo) can improve the bonding force between Ni and Al atoms, leading to the decrease of diffusion rate; (4) grain boundaries in NiCrAl coating can promote elements diffusion, but there are no grain boundaries in the single-crystalline substrate. 3.3. Diffusion resistance of the diffusion barrier
Fig. 6. Standard Gibbs free energy for Al2O3, Cr2O3, NiO and [Zr] as function of temperature.
Fig. 9 shows the cross-sectional morphology and element profiles
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along the thickness of the N5/DBC/NiCrAl system that exposed at 1000 °C for 200 h in the atmosphere. It can be found that the α-Cr phase has not formed in the NiCrAl coating (Fig. 10a). The IDZ, SRZ, and voids were not formed compared with the N5/NiCrAl system, and a good adherence of the interface was observed. The elements profiles along the thickness of the N5/DBC/NiCrAl system were measured as shown in Fig. 10b–j. The selected areas for EDS are labeled with number (points 1 to 7 in Fig. 10a). The O and Zr were only detected in the sandwiched diffusion barrier. The diffusion of Ni and Cr cannot be observed from Fig. 10d and e. A small amount of Co was found in the α-Al2O3 layer adjacent to substrate side, which may be attribute that the alumina diffusion barrier was not dense in the initial formation stage, so that the Co was diffused to the α-Al2O3 layer. From Fig. 10h–j, it can be seen that the diffusion of the refractory elements W, Ta, and Re was blocked by the DBC. The aluminum as an importantly diffused element, its contents from point 1 to 7 were 3.12, 2.04, 31.03, 8.85, 50.17, 0 and 0 wt%, respectively. It can be noted that the depletion of Al occurred in the NiCrAl coating close to the DBC, which would suppress the growth of the α-Al2O3 in the diffusion barrier coating. In the N5/NiCrAl system, the elements interdiffusion were caused due to the inward diffusion of the alloying elements from coating to substrate and the outward diffusion of the refractory elements from substrate to NiCrAl coating. The contents of the refractory elements were relatively low in the substrate, which resulted in the low concentration gradient along the diffusion direction. As is well known, the diffusion coefficient of the refractory elements is normally lower than that of the relatively low melting metal elements. Therefore, the contents and diffusion coefficients of the alloying elements in NiCrAl coating are higher than that of the refractory elements in the substrate. These results suggest that the diffusion fluxes of atoms from NiCrAl coating to substrate were higher than that from substrate to coating, leading to the formations of the IDZ, SRZ and Kirkendall porosities. Additionally, the diffusion of Al was the main factor of affecting the performance of the coating and substrate during the elements interdiffusion. According to the first principal of Fick's law:
Fig. 7. XPS spectra of the Al element (a) and Zr element (b) in the interface reaction zone between ZrO2/NiCrAl coating under as-deposited state.
formation of the IDZ was resulted from the interdiffusion of Al, Cr, and Ni at the interface of substrate and NiCrAl coating, which can promote the precipitation of refractory elements (e.g. Re, W, and Ta) to form SRZ zone and needle-like TCP phase in substrate. In addition, some dark phases were precipitated in the NiCrAl coating, and some Kirkendall voids were observed at the IDZ/NiCrAl interface. Elemental distribution along thickness direction for the N5/NiCrAl system was measured by EDS (Fig. 9b–h). The trace to measure the elemental profile is a row of points (numbered 1–6) shown in Fig. 9a. Variation of the Ni, Cr, Al, and Co elements in the IDZ can be ascribed to elemental interdiffusion. The high content of Cr (87.4 in wt%) in the dark phase (point 4) suggests that the phase should be α-Cr. A certain amount of W (~2.7%) was detected in the α-Cr phase (Fig. 9e), implying that W might have played some role in the formation of α-Cr. The existence of Co (Fig. 9f) and W elements in the NiCrAl coating indicated that these two elements diffused from the substrate across the interface into the NiCrAl coating layer. Fig. 10 shows the cross-sectional morphology and element profiles
J=
D
d dx
(3)
where J is the diffusion flux, D is the diffusion coefficient, ρ is the mass concentration of the diffused substance, and dρ/dx is the concentration gradient along the thickness direction. The diffusion flux of Al is proportional to the diffusion coefficient and the mass concentration, but inversely proportional to the diffusion distance that can be regarded as the thickness of the diffusion barrier. For the N5/ZrO2/NiCrAl system, the initial single ZrO2 precursor layer will transform into a sandwich structure when Al reacts with ZrO2 continuously to form Al2O3. That led to the increase of diffusion barrier thickness, i.e., the diffusion distance. Simultaneously, the alloying elements have the lower diffusion rate in α-Al2O3 than that in the DBC and NiCrAl coating [17]. Therefore, the interdiffusion of the metal elements would become much more difficult, and the formation of IDZ, SRZ and voids were suppressed due to the existence of the DBC.
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Fig. 8. Cross-sectional morphology of diffusion barrier exposed at 1000 °C for 5 h (a), 100 h (b), 200 h (c) and the variation of average thickness with the oxidation time (d).
3.4. Surface morphologies of the heat-treated samples
interdiffusion of the two systems were studied. Conclusions can be drawn as follows:
The surface morphologies of NiCrAl coatings in the N5/NiCrAl and N5/DBC/NiCrAl systems were shown in Fig. 11. The spallation of the scales can be observed in the two systems. The XRD patterns of the NiCrAl surfaces in the two systems after 200 h exposure at 1000 °C in atmosphere were shown in Fig. 12. As can be seen, the surfaces of the two coatings were composed of γ-Ni and α-Al2O3. The results revealed that the original β-NiAl and the γ´-Ni3Al phases in Fig. 3 have transformed into the γ-Ni phase because of the formation of the α-Al2O3. The generation of α-Al2O3 scale in surfaces could improve the oxidation resistance of the systems.
(1) The interdiffusion zone (IDZ), secondary reaction zone (SRZ) and Kirkendall voids were formed in the N5/NiCrAl system, due to the interdiffusion of elements after thermal exposure at 1000 °C for 200 h. (2) For the N5/ZrO2/NiCrAl system, the ZrO2 precursor layer transformed into a sandwich structural DBC (α-Al2O3/rich-Zr/α-Al2O3) after thermal exposure at 1000 °C, and the average thickness of each α-Al2O3 layer increased with the increase of thermal exposure time. (3) The DBC could suppress the interdiffusion of the alloying elements compared to the system without the diffusion barrier coating.
4. Conclusions The effect of thermal exposure on the structure evolution of the N5/ NiCrAl and N5/ZrO2/NiCrAl systems was analyzed. The elements
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Fig. 9. Cross-sectional morphology (a) and elements profiles (b–h) along the thickness of the N5/NiCrAl system after 200 h exposure at 1000 °C in atmosphere.
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Fig. 10. Cross-sectional morphologies (a) and element profiles (b-j) along the thickness of the N5/DBC/NiCrAl system after 200 h exposure at 1000 °C in atmosphere.
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Acknowledgments This work was supported by National Key Basic Research Program of China (973 Program, No. 2012CB625102). References [1] Q. Lu, Y.L. Pei, S.S. Li, et al., Role of volatilization of molybdenum oxides during the cyclic oxidation of high-Mo containing Ni-based single crystal superalloys, Corros. Sci. 129 (2017) 192–204. [2] N.D. Souza, J. Kelleher, C.L. Qiu, et al., The role of stress relaxation and creep during high temperature deformation in Ni-base single crystal superalloys - implications to strain build-up during directional solidification, Acta Mater. 106 (2016) 322–332. [3] R. Farangis, Z.M. Li, S. Zaefferer, et al., On the origin of creep dislocations in a Nibase, single-crystal superalloy: an ECCI, EBSD, and dislocation dynamics-based study, Acta Mater. 109 (2016) 151–161. [4] M.J. Pomeroy, Coatings for gas turbine materials and long term stability issues, Mater. Des. 26 (2005) 223–231. [5] M. Elsaß, M. Frommherz, A. Scholz, et al., Interdiffusion in MCrAlY coated nickelbase superalloys, Surf. Coat. Technol. 307 (2016) 565–573. [6] H.R. Yao, Z.B. Bao, M.L. Shen, et al., A magnetron sputtered microcrystalline β-NiAl coating for SC superalloys. Part II. Effects of a NiCrO diffusion barrier on oxidation behavior at 1100 °C, Appl. Surf. Sci. 407 (2017) 485–494. [7] A. Sato, H. Harada, K. Kawagishi, Development of a new bond coat “EQ coating” system, Metall. Mater. Trans. A 37A (2006) 789–790. [8] K. Kawagishi, A. Sato, H. Harada, A concept for the EQ coating system for nickelbased superalloys, JOM-USL 60 (2008) 31–35. [9] K. Kawagishi, H. Harada, Akihiro Sato, et al., EQ coating: a new concept for SRZfree coating systems, Superalloys (2008) 761–768. [10] Z. Bai, D. Li, H. Peng, et al., Suppressing the formation of SRZ in a Ni-based single crystal superalloy by RuNiAl diffusion barrier, Prog. Nat. Sci. 22 (2012) 146–152. [11] T. Narita, F. Lang, K. Thosin, et al., Oxidation behavior of nickel-base single-crystal superalloy with rhenium-base diffusion barrier coating system at 1,423 K in air, Oxid. Met. 68 (2007) 343–363. [12] E. Cavaletti, S. Naveos, S. Mercier, et al., Ni–W diffusion barrier: its influence on the oxidation behaviour of a β-(Ni, Pt) Al coated fourth generation nickel-base superalloy, Surf. Coat. Technol. 204 (2009) 761–765. [13] L.J. Zhu, S.L. Zhu, F.H. Wang, Preparation and oxidation behavior of nanocrystalline Ni+CrAlYSiN composite coating with AlN diffusion barrier on Ni-based superalloy K417, Corros. Sci. 60 (2012) 265–274. [14] S.H. Hosseini, S. Mirdamadi, S. Rastegari, Investigating efficiency of α-Al2O3 diffusion barrier layer in oxidation of EB-PVD NiCrAlY coatings, Surf. Eng. 31 (2015) 146–155. [15] P. Ren, S.L. Zhu, F.H. Wang, Microstructural stability of AlN diffusion barrier for nanocompositeNi + CrAlYSiHfN coating on single crystal superalloy at high temperatures, Appl. Surf. Sci. 359 (2015) 420–425. [16] Y.X. Xu, X.T. Luo, C.X. Li, Effect of annealing conditions on in situ formation of Cr2O3 diffusion barrier, Surf. Eng. 33 (2017) 210–216. [17] Q.M. Wang, Y.N. Wu, M.H. Guo, et al., Ion-plated Al-O-N and Cr-O-N films on Nibase superalloys as diffusion barriers, Surf. Coat. Technol. 197 (2005) 68–76. [18] W.Z. Li, Q.M. Wang, Z.B. Bao, et al., Microstructural evolution of the NiCrAlY/CrON duplex coating system and its influence on mechanical properties, Mater. Sci. Eng. A 498 (2008) 487–494. [19] C.A. Guo, W. Wang, Y.X. Cheng, et al., Yttria partially stabilised zirconia as diffusion barrier between NiCrAlY and Ni-base single crystal René N5 superalloy, Corros. Sci. 94 (2015) 122–128. [20] S. Merchant, M.R. Notis, A review – constitution of the Al–Cr–Ni system, Mater. Sci. Eng. 66 (1984) 47–60. [21] I. Barin, O. Knacke, O. Kubaschewski, Thermochemical Properties of Inorganic Substances, Springer-Verlag, Berlin, 1991. [22] S. Yu, T. Yu C Wang, et al., Self-diffusion in the intermetallic compounds NiAl and Ni3Al-An embedded atom method study, Physica B 396 (2007) 138–144.
Fig. 11. Surface morphologies of the NiCrAl after 200 h exposure at 1000 °C in atmosphere: N5/NiCrAl system (a) and N5/DBC/NiCrAl system (b).
Fig. 12. XRD patterns of the NiCrAl surface after 200 h exposure at 1000 °C in air.
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