A magnetron sputtered microcrystalline β-NiAl coating for SC superalloys. Part I. Characterization and comparison of isothermal oxidation behavior at 1100 °C with a NiCrAlY coating

Applied Surface Science 324 (2015) 1–12

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A magnetron sputtered microcrystalline ␤-NiAl coating for SC superalloys. Part I. Characterization and comparison of isothermal oxidation behavior at 1100 ◦ C with a NiCrAlY coating Shaojun Hou a,b , Shenglong Zhu a,∗ , Tao Zhang a,b , Fuhui Wang a,b a

Institute of Metal research, Chinese Academy of Sciences, 62 Wencui Road, Shenyang 110016, China Corrosion and Protection Laboratory, Key Laboratory of Superlight Materials and Surface Technology (Harbin Engineering University), Ministry of Education, Nantong ST 145, Harbin 150001, China b

a r t i c l e

i n f o

Article history: Received 19 September 2014 Received in revised form 19 October 2014 Accepted 19 October 2014 Available online 28 October 2014 Keywords: Microcrystalline ␤-NiAl Magnetron sputtering High temperature oxidation Inter-diffusion Scale adherence

a b s t r a c t A microcrystalline ␤-NiAl coating was prepared on a single-crystal (SC) superalloy substrate via magnetron sputtering and subsequent vacuum annealing. The grain sizes of the coating ranged from about 300 nm to 1 ␮m. A reference NiCrAlY coating, which was mainly comprised of ␥ -Ni3 Al and ␣-Cr, was prepared by means of vacuum arc evaporation (VAE). Isothermal oxidation tests were carried out at 1100 ◦ C in air for 50 h. Both coatings formed thin and adherent ␣-Al2 O3 scales during tests, while the oxide scales on the bare superalloy primarily consisted of spinel (Ni,Co)Al2 O4 with underlying ␣-Al2 O3 scale. The parabolic rate constant of the NiAl-coated specimens was about one order of magnitude lower than that of the NiCrAlY coated specimens. After oxidation tests, only a small amount of ␥ phase was detected at some columnar boundaries of the ␤-NiAl coating, and about 2/3 parts of the NiCrAlY coating transformed into ␥ phase which resolved the ␣-Cr precipitations, while an Al-depleted zone in thickness of about 10 ␮m formed beneath the TGO of the bare superalloy. Inter-diffusion zones and secondary reaction zones were observed on the specimens coated by either ␤-NiAl or NiCrAlY. The oxidation mechanism and microstructure evolvement of the specimens during high temperature exposures were discussed. © 2014 Elsevier B.V. All rights reserved.

1. Introduction To satisfy the demands of higher efficiency of gas turbine engines, the inlet gas temperatures are increasing progressively. Therefore, it is essential to develop single crystal (SC) superalloys and new thermal barrier coating (TBC) systems capable of withstanding higher surface temperatures. A TBC system typically consists of an yttria-partially-stabilized zirconia top coat providing thermal insulation and an Al-rich oxidation resistant metallic bond coat (BC). During high temperature exposures, a thermally grown oxide (TGO) layer forms at the interface between the bond coat and the top coat as a consequence of the reaction between the oxygen penetrating through top coat and the metallic elements diffusing from the bond coat [1–5]. It is commonly agreed that the formation and growth behavior of TGO during service is one of the most important issues responsible for the failure of TBCs [2]. An ideal TGO should be slow growing, continuous, and adherent to BC. As ␣-Al2 O3 has extremely low self

∗ Corresponding author. Tel.: +86 024 23904856; fax: +86 024 23893624. E-mail address: [email protected] (S. Zhu). http://dx.doi.org/10.1016/j.apsusc.2014.10.106 0169-4332/© 2014 Elsevier B.V. All rights reserved.

and impurity diffusion rates, and its chemical and thermal stability is relatively high, ␣-Al2 O3 is preferred to other oxides such as NiO, Cr2 O3 , Ni(Al,Cr)2 O4 , and ␪-Al2 O3 [4–7]. To form and maintain the growth of ␣-Al2 O3 scale at high temperatures, a high Al content is demanded for a good BC. MCrAlY (M = Ni, Co, or their combination) is one of the most common BC materials considering its balanced oxidation and hot corrosion resistance [3,8]. However, at very high temperatures, e.g. 1100 ◦ C or higher, Al in MCrAlY may be depleted too fast, which leads to forming non-alumina oxides and then premature failure of TBCs [8,9]. Besides, faster TGO growth rates may accelerate the spallation failure of TBCs because the stresses at the TGO/BC interface increase with the thickness of the TGO layer [10–12]. So, ␤-NiAl as a BC material is more attractive at higher temperatures, e.g. 1050 ◦ C at which hot corrosion is ignorable and TGO growth rates is high, because higher Al contents in NiAl coating provide longer life-time than NiCrAlY coating. However, the ␤-NiAl coatings prepared by conventional diffusion method or by chemical vapor deposition show poor cyclic oxidation resistance due to voids formation at the TGO/coating interface and surface rumpling which lead to scale spallation [13–15]. Numerous approaches to improve the cyclic oxidation performance of NiAl have been carried out, e.g.

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Table 1 Nominal composition of the SC superalloy substrate. Chemical composition (wt.%) Element

Ni

Co

Cr

Al

Ta

W

Re

Mo

Hf

C

Y

B

SC substrate

Bal

7.5

7.5

6.2

6.5

5.0

3.0

1.5

0.15

0.05

0.01

0.004

Pt alloying to form NiPtAl, reactive elements (RE) (e.g. Hf, Y, Zr, etc.) or their oxides modifications, and grain refinement (nanocrystallization or microcrystallization) [14–19]. Recently Guo et al. [20,21] indicated that Hf/Zr and Y/La co-doping can lower the oxidation rate of ␤-NiAl at 1200 ◦ C further than single-doped alloys, by mechanisms of co-segregation of the RE ions on oxide grain boundaries and acting as an ionic cluster which has stronger interaction with Al ions than the individual ions. However co-doping of Hf/Dy as well as Hf/La in ␤-NiAl did not have the same synergistic effect on the growth of alumina scale because of their relatively smaller ‘effective radius’. Hou [22] found that Pt in NiAl eliminated S segregation at the TGO/alloy interface and reduced the extent of interfacial pore formation thus improved the TGO adhesion. Svensson [16] pointed out that the beneficial effect of Pt is also attributed to the enhancement of contact areas between TGO and metal in addition to the enhanced interface bondings. Lou et al. [23] studied the oxidation behavior of sputtered micrograined K38G coating and implied that microcrystallization not only reduced the critical aluminum content necessary to form alumina on the surface of K38G superalloy but also improved the scale adhesion of the superalloy. Chen [24] also reported that sputtered nanocrystalline Ni–8Cr–3.5Al coating showed better oxidation resistance than the corresponding cast alloy and it was closely related to the rapid diffusion of Al through grain boundaries in the nanocrystalline coating. A few reports are focused on the oxidation behavior of nano- or micro-grained NiAl coatings on superalloys at 900 ◦ C or 1000 ◦ C [19,25]. However the oxidation behavior of nano- or micro-grained NiAl coatings on SC superalloys at very high temperatures is still unclear up to date. In the present work, a microcrystalline ␤-NiAl coating on a SC superalloy substrate was deposited by magnetron sputtering from multi-targets of Ni3 Al and Al. ␤-NiAl is brittle and prone to crack during manufacture and coating deposition, especially for large size sputtering targets. The isothermal oxidation tests of the ␤-NiAl coating was conducted at 1100 ◦ C in air. The oxidation performance and microstructure evolvements of the coating were studied and discussed in detail in this article, and compared with that of a conventional NiCrAlY coating prepared by vacuum arc evaporation (VAE) as well as the bare SC superalloy substrate. The effect of active diffusion barrier on the isothermal and cyclic oxidation behavior of the microcrystalline NiAl coating will be reported in detail in Part II. 2. Experimental procedures 2.1. Preparation of coatings A Ni-based single-crystal superalloy was used as the substrate, whose nominal composition (wt.%) is shown in Table 1. Each specimen measuring approximately 16 mm in diameter and 1.5 mm in height was cut from SC bars. A hole with a diameter of 1.5 mm near the edge of each specimen was made by electrical-spark cutting, for the purpose of hanging and rotating the specimen during vacuum deposition so as to allow all faces to be uniformly coated. Then the specimens were grounded down to 1000-grit emery paper, degreased in acetone and followed by 10 s etching in CuSO4 solution to ensure coating adhesion. All the specimens were ultrasonically cleaned in alcohol and dried in hot air immediately before coating deposition. For sputter deposition, an Al target (99.99% pure) and duplex Ni3 Al targets were used. Dimensions of the targets

Table 2 Sputtering parameters for NiAl coating. Sputtering conditions

Parameters

Targets Sputtering current (A) Frequency (kHz) Duty cycles Base pressure (Pa) Ar pressure (Pa) Substrate temperature (◦ C) Deposition time (h)

Ni3Al 7.0 40 80%

Al 3.3 – – 6.0 × 10−4 0.35–0.40 180–200 5.5

were 380 mm × 128 mm × 8 mm. Two Ni3 Al targets faced to each other, while the Al target was perpendicular to them. The Al target was connected to a DC sputtering power supply while two identical Ni3 Al targets were connected to a mid-frequency square wave power supply. All the power supplies worked in constant current mode. For the Ni3 Al targets, one target worked as anode and another as cathode in each first half-cycle, and swapped their anode/cathode functionality in the following half-cycle. The process parameters used in sputtering are shown in Table 2. For comparison, some specimens were coated by NiCrAlY coating using vacuum arc evaporation method from a target Ni–27Cr–11Al–0.5Y (wt.%) and detailed deposition parameters have been reported elsewhere [26]. After deposition all the specimens were annealed at 1000 ◦ C for 1 h in a quartz tube furnace in vacuum about 6.0 × 10−4 Pa. 2.2. High temperature oxidation tests Isothermal oxidation of microcrystalline ␤-NiAl coated, NiCrAlY coated and uncoated SC specimens was conducted at 1100 ◦ C for 50 h in a vertical microbalance (Thermax 700 thermal microbalance, Thermo Cahn, USA) with a sensitivity of 10−5 g. After oxidation, each specimen was cooled to room temperature in the chamber. 2.3. Characterization The surface and cross-sectional morphologies and chemical compositions of the coatings were characterized using scanning electron microscopy (SEM, Inspect F 50, FEI Co., Hillsboro, OR) equipped with energy dispersive X-ray spectrometer (EDS, X-Max, Oxford instruments Co., Oxford, UK) and electron probe microanalysis (EPMA-1610, Shimadzu, Kyoto, Japan). X-ray diffraction (XRD, X’Pert PRO, PANalytical Co., Almelo, Holland, Cu Ka radiation at 40 kV) and transmission electron microscopy (TEM, JEM-2100F, JEOL, Tokyo, Japan) were used for phase identification. TEM was also used to investigate the microstructure details of the coatings before and after oxidation. For cross-sectional observation, an electrolessnickel layer was applied on the surface of the specimens and then mounted in epoxy resin before grounding and polishing. 3. Results 3.1. Fabrication and characterization of the coatings Fig. 1 gives the XRD patterns of as-sputtered and as-annealed NiAl coatings. The as-deposited coating consists of ␤-NiAl and

S. Hou et al. / Applied Surface Science 324 (2015) 1–12

Fig. 1. XRD patterns of as-sputtered and as-annealed NiAl coatings.

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minor Ni2 Al3 and Ni3 Al4 phases. The peaks of Ni2 Al3 and Ni3 Al4 phases are a little broadened probably because of their small grain sizes. The average chemical composition of as-deposited coating is Ni–52Al (at.%) according to EPMA analysis. The concentration of Al in ␤-NiAl can be thermodynamically stable as high as 55 at.% as referred to the Ni–Al equilibrium phase diagram [27]. Minor Ni2 Al3 and Ni3 Al4 phases were detected in the as-sputtered coating, probably because the coating was deposited from multi-targets and therefore there is non-uniformity of chemical compositions in the coating. After annealing for 1 h at 1000 ◦ C, Ni2 Al3 and Ni3 Al4 phases vanished and a single phase ␤-NiAl coating with strong (1 1 1) and (2 1 1) preferred orientations was obtained. It is noted that the main diffraction peaks of ␤-NiAl, (1 1 1) and (2 1 1), shifted slightly to lower diffraction angles after annealing. This results from lattice expansion during phase transforming from Al-rich Ni0.9 Al1.1 to Nirich Ni1.1 Al0.9 due to coating/substrate inter-diffusion which will be discussed in the following paragraph. Surface and cross-sectional morphologies of as-sputtered and as-annealed NiAl coatings are presented in Fig. 2. The as-deposited NiAl coating shows typical columnar microstructures [19]. The column sizes are less than 3 ␮m as shown in the inserted high magnification figure in Fig. 2a. The coating with a thickness of about 31 ␮m is adherent to the underlying substrate (Fig. 2c). The white

Fig. 2. Surface and cross-sectional morphologies of as-sputtered (a and c) and as-annealed NiAl coatings (b and d).

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arrows in Fig. 2c present wide columnar boundaries which are commonly observed in thick sputtered coatings [28]. The boundary between the as-sputtered coating and the substrate was clear, indicating there was nearly no coating/substrate inter-diffusion during sputter deposition owing to the relatively low deposition temperature (180–200 ◦ C). The vacuum annealing led to little changes in the surface morphology as shown in Fig. 2b. The columnar microstructure, however, disappeared in the cross-sectional SEM image in Fig. 2d. Besides, annealing at 1000 ◦ C resulted in an inter-diffusion zone (IDZ) between the coating and substrate. The IDZ is about 5 ␮m thick, and rich in refractory metals, e.g. Mo, W and Re according to the EDX results. During high temperature annealing, the chemical composition gap between the coating and substrate provided the primary driving force for coating/substrate inter-diffusion which enhanced the adhesion between the coating and substrate [29]. The white arrows in Fig. 2d indicate a few small Kirkendall voids along the coating side above IDZ. They are the evidences that inward diffusion is dominant. The average Al concentration in the as-annealed coating slightly decreased to about 48 at.% from 52 at.%, according to the analyses of EDX and EPMA. As a result, the Al rich sputtered coating transformed to a slightly Ni rich single ␤-NiAl coating. In order to identify the grain size of as-annealed coating, planar view TEM observation was conducted as shown in Fig. 3. The

Fig. 3. Planar view TEM micrograph for as-annealed ␤-NiAl coating.

grain sizes of the coating are varied ranging from 300 nm to about 1 ␮m, and many high density dislocations are clearly shown in some grains. The surface and cross-sectional morphologies as well as XRD pattern of as-annealed NiCrAlY coating are displayed in Fig. 4. The

Fig. 4. Surface (a), cross-sectional morphologies (b) and XRD pattern of as-annealed VAE NiCrAlY coating (c).

S. Hou et al. / Applied Surface Science 324 (2015) 1–12

NiCrAlY coating was 25–31 ␮m thick, consisted mainly of ␥ phases with uniformly distributed ␣-Cr precipitations (darker particles in Fig. 4b). And minor ␤-NiAl phase was identified by XRD combined with TEM (not shown here). It is notable that two extremely weak peaks of Al2 Y4 O9 phase were observed in the XRD pattern. The formation of Al2 Y4 O9 phase may result from oxidation of Al and Y under extreme low oxygen pressure during the annealing process by: 4Al + 8Y + 9O2 = 2Al2 Y 4 O9 .

(1)

The average chemical composition of the NiCrAlY coating has been determined by EDS to be 61.5Ni–0.8Co–27.7Cr–10.3Al in wt.%. No inter-diffusion zone between the coating and substrate can be clearly distinguished, though evidences for coating/substrate inter-diffusion were found. Minor Co was detected in the coating, and EDS results reveal that the Al content right beneath the coating/substrate interface was 7.5 wt.%, which was 6.2 wt.% before annealing. 3.2. Isothermal oxidation at 1100 ◦ C 3.2.1. Oxidation kinetics The oxidation kinetics curves of microcrystalline ␤-NiAl coating and conventional NiCrAlY coating as well as bare superalloy for 50 h isothermal oxidation at 1100 ◦ C are illustrated in Fig. 5a. It is noted that the substrate has the lowest mass gain. However

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Table 3 Parabolic rate constants at 1100 ◦ C and average scale thickness of ␤-NiAl coating and NiCrAlY coating after 50 h of oxidation.

␤-NiAl coating NiCrAlY coating

K1 (g2 /cm4 s)

K2 (g2 /cm4 s)

Average scale thickness (␮m)

3.2 × 10−12 2.8 × 10−12

6.3 × 10−14 4.6 × 10−13

1.74 2.91

fluctuations of the curve are obvious, indicating that oxide spallation and rescale occurred during the isothermal-oxidation process. The curve fittings in Fig. 5b and c demonstrate that the oxidation rates of both coatings follow two-stage parabolic laws. The parabolic rate constants and the average thickness of the oxide scales calculated from the cross-sectional micrographs are listed in Table 3. For ␤-NiAl coating, it exhibits a fast scaling rate in the first ∼1 h (the first stage), and then shows a slow steady-state scaling rate (the second stage) and the total mass gain is 0.19 mg/cm2 . The scaling rate in the second stage of NiCrAlY coating was almost one order of magnitude higher than that of NiAl coating, and the total mass gain is 0.37 mg/cm2 , nearly twice of the ␤-NiAl coating. The first stages of both coatings had almost the same parabolic constants. In fact, they correspond to the growth of meta-stable ␪Al2 O3 scales as indicated by XRD analysis which will be shown later in Fig. 9. And the growth rates in the second stage dropped because ␪-Al2 O3 growth gave place to ␣-Al2 O3 scale growth. Extremely

Fig. 5. Isothermal oxidation kinetics of the coatings and substrate at 1100 ◦ C in air (a) and curve fittings of microcrystalline ␤-NiAl coating (b) compared with that of conventional VAE NiCrAlY coating (c).

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Fig. 6. XRD patterns of the NiCrAlY coating (a), microcrystalline ␤-NiAl coating (b) and SC superalloy substrate (c) after 50 h of oxidation at 1100 ◦ C in air.

weak XRD peaks of Al2 Y4 O9 were observed in as-annealed NiCrAlY coating. So, Y may play important roles in higher oxidation rates for the NiCrAlY coating. 3.2.2. Oxidation products and morphologies after oxidation Fig. 6 gives the XRD patterns of the specimens after 50 h of oxidation at 1100 ◦ C. Only ␣-Al2 O3 was detected on the ␤-NiAl coating and NiCrAlY coating, while mixed oxides including ␣-Al2 O3 , (Ni,Co)Al2 O4 and NiTa2 O6 were detected on the SC substrate. In addition, ␥ -Ni3 Al phase could be hardly found in the XRD patterns of the oxidized microcrystalline NiAl coating. The ␣-Cr was not able to be detected in the NiCrAlY coating. Most surfaces of the oxidized bare SC superalloy are smooth, while some large ridges can be seen as indicated by arrows in Fig. 7a. The oxide ridges are Ta-rich oxides intruded and extrude the substrate surface (Fig. 7d), and an alumina layer beneath them extended to other parts of the oxide scale. Combining the local area EDS analysis and XRD results, one may conclude that the Ta-rich ridges are NiTa2 O6 . The main part of the oxide scale consisted of an inner layer of alumina and an outer layer of spinel (Ni,Co)Al2 O4 according to EDS analysis. An Al-depleted zone beneath TGO is about 10 ␮m thick as shown in Fig. 7d, which resulted from Al consumption for the growth of alumina-rich TGO. A uniform TGO scale formed and was adherent to the NiCrAlY coating (Fig. 7e). It is wavier than that on NiAl coating, because the as-prepared NiCrAlY coating had coarser surface than sputtered NiAl coating. Fig. 8 shows more details of the oxide formed on NiCrAlY coating. Several white inclusions were inserted into the alumina scale as indicated by white arrows in Fig. 8a. They are identified as Y-rich oxides as seen in the EDS spectrum in Fig. 8b. Considering weak Al2 Y4 O9 peaks shown in Fig. 4c, one may conclude that these Y-rich oxides are spinel Al2 Y4 O9 too. The higher rates of TGO growth on the NiCrAlY coating may be attributed to the formation of Al2 Y4 O9 . Though ␣-Cr diffraction peaks vanished, a few Cr-rich particles were found in the coating as indicated by black arrows. In comparison with the cross-sectional morphology of the as-annealed NiCrAlY coating (Fig. 4b), the oxidized NiCrAlY coating shows some new features (Fig. 7e). Large blocks of brighter phase can be observed within the lower part of the coating. EPMA analysis, which will be shown in Fig. 11, indicates that they were Al-rich phases and probably ␥ -Ni3 Al, i.e. the NiCrAlY coating without degradation. Above them, there was an Al-depleted zone due to Al consumption for the TGO growth. And beneath them, a thin and continuous brighter phase zone appeared

at the original coating/substrate interface. EPMA analysis in Fig. 11 indicates that it was a Cr-rich zone. Beneath the Cr-rich zone, there was an IDZ, an almost uniform ␥ zone. In the SRZ, Re and W rich needle-shaped phases were observed. Fig. 9 gives the surface morphologies and glancing incidence X-ray diffraction patterns (GI-XRD) of the NiAl-coated specimens after oxidation for 1 h and 50 h, respectively. Fig. 9a indicates that ␪-Al2 O3 and ␣-Al2 O3 coexisted, which is conformed by the typical needle or blade like morphologies shown in Fig. 9b. After 50 h of oxidation, no ␪-Al2 O3 could be found, suggesting the ␪-␣-Al2 O3 phase transformation has completed. In fact, the initial rapid oxidation of ␤-NiAl coating was commonly thought to be a consequence of the formation of fast-growing meta-stable ␪-Al2 O3 [13]. Once the growth of ␣-Al2 O3 dominates, the growth rates will drop down because ␣-Al2 O3 has parabolic rate constants about two orders of magnitude smaller than those of ␪-Al2 O3 at the same temperatures [13,30]. One can notice that, by comparing Fig. 7c with Fig. 2b, the surface morphology of the oxidized NiAl coating is almost identical to that of the as-annealed one. The cross-sectional view (Fig. 7f) reveals a continuous and adherent TGO which is ∼40% thinner than that on the NiCrAlY coating. The inserted high magnification picture in Fig. 7f displays a thin sting shaped ␥ -Ni3 Al phase extended from the TGO/coating interface into the ␤-NiAl coating because of ␤/␥ phase transformation due to the Al consumption for alumina growth. It is noticed that the distribution of ␥ -Ni3 Al phases in the NiAl coating were probably along the original columnar grain boundaries which acted as fast paths for outward diffusion of Al. An inter-diffusion zone (IDZ) and a secondary reaction zone (SRZ) are more obvious and slightly thicker in comparison with those on the NiCrAlY coated specimens, apparently because the Al contents in the former coating were much higher. EDS analysis indicated that the matrix phase in IDZ was ␤-NiAl and that in SRZ was ␥ -Ni3 Al, which exhibited a step-down gradient of Al contents. Besides, many large ␥ grains existed in IDZ. TCP precipitated as coarse particles in IDZ and as fine rods in SRZ. The density of TCP phases in SRZ of NiAl-coated specimens was at least three times of that of NiCrAlY coated specimens. 3.2.3. Characterization by EPMA mapping Cross-sectional elemental mappings of ␤-NiAl coating and NiCrAlY coating after 50 h of isothermal oxidation at 1100 ◦ C in air were obtained through EPMA method, as shown in Figs. 10 and 11, respectively. The O and Al maps in Fig. 10 indicate that a continuous alumina scale formed on the ␤-NiAl coating and some fine alumina micro-pegs penetrated into the coating. Besides, there were a great amount of Co and Cr along with a few W, Mo and Re, which must come from the substrate. Beneath the coating, an IDZ (Zone I) and a SRZ (Zone II) exhibit their distinctive chemical compositions. The IDZ comprises at least two phases, one had quite strong Al signals which was comparable to that of the ␤-NiAl coating, and another was rich in refractory elements of Re, W, Mo and Cr. In fact, the IDZ was composed of ␤ matrix with scattered ␥ and TCP, according to previous analysis. The average contents of every alloying element in the SRZ were between those in the IDZ and the un-affected substrate. Fig. 11 shows the elemental mapping of NiCrAlY coating after 50 h of isothermal oxidation at 1100 ◦ C in air. A continuous alumina scale along with a few pegs was observed. Y was mainly located within the alumina scale and at the original coating/substrate interface. A few Ta were detected within the coating. Segregation of Cr was observed within the coating and at the coating/substrate interface. It is noted that the Cr rich region was also abundant in Re, Mo and W, and was depleted in Ni and Al. The particles rich in Cr and refractory elements formed an almost continuous layer at the coating/substrate interface. They are probably the cause of the nearby

S. Hou et al. / Applied Surface Science 324 (2015) 1–12

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Fig. 7. Surface and cross-sectional morphologies of SC superalloy substrate (a and d), NiCrAlY coating (b and e) and ␤-NiAl coating (c and f) after 50 h of isothermal oxidation at 1100 ◦ C in air.

Cr-depleted zone in the coating. Beneath them, there was a very thin layer rich in Al and Y, which was followed by another very thin Cr-rich layer and then a relatively thicker transition zone. It is obvious that the bright needle phases beneath the transition zone were rich in Re and depleted in Cr. 4. Discussion 4.1. Growth rates and adherence of TGO During high temperature exposure, the formation and maintenance of a continuous, adherent and slow growing ␣-Al2 O3 scale provides excellent oxidation resistance for the superalloys and metallic coatings. To form a continuous external ␣-Al2 O3 scale, the concentration of Al in superalloys and coatings must exceed

a critical value NAl according to Wanger’s theory [31]. The value NAl is defined as NAl =

 g ∗ N D V 1/2 O O M 2DAl VOX

(2)

where DO NO is the oxygen permeability in the alloy; DAl is the diffusion coefficient of Al in the alloy; g* is a factor determined by the volume fraction of the oxide required for the transition from internal to external oxide formation, and VM and VOX are the molar volumes of the alloy and oxide, respectively. In this work, both the ␤-NiAl coating and the NiCrAlY coating formed exclusive ␣-Al2 O3 scales while the uncoated SC superalloy formed mixed oxides including ␣-Al2 O3 , (Ni,Co)Al2 O4 and NiTa2 O6 . The SC superalloy used in this work has 13.8 at.% Al and 8.1 at.% Cr. At least 31 at.% Al is required for Ni–Al binary

Fig. 8. Cross-sectional morphology of oxide formed on NiCrAlY coating after 50 h of oxidation at 1100 ◦ C (a) and the corresponding EDS spectrum of P1 (b).

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Fig. 9. GI-XRD patterns (a) and surface morphologies of ␤-NiAl coating after 1 h (b) and 50 h of oxidation at 1100 ◦ C, respectively.

alloys to form and maintain an external Al2 O3 at temperatures from ∼900 ◦ C to ∼1150 ◦ C, otherwise a thick scale containing both NiO and spinel NiAl2 O4 will form above a discontinuous alumina layer [32]. The addition of Cr to Ni–Al alloys results in a remarkable synergistic effect to enhance the formation of external alumina scale. For example, a model alloy Ni–10.5Cr–10Al (at.%) and a commercial superalloy DZ125 (Ni–10.2Cr–11.2Al–1.2Ti–10.0Co–1.2Ta–1.2Mo–2.3W–0.0055B, at.%) are able to form external alumina scale. Ni–11Cr–8Al (at.%), however, formed a Cr2 O3 external scale with internal alumina layer [33]. If drawing a line from point A(31Al,0Cr) to point B(10Al,10.5Cr), one may note that both the SC superalloy used in this work and a ternary alloy Ni–11Cr–8Al (at.%) are located just below the line: 2NCr + NAl = 31

(3)

where NCr and NAl are the contents in atom percent of Cr and Al in the alloys, respectively. So Eq. (3) plus a minimal Al requirement, e.g. 10 at.% as inferred from Ni–10.5Cr–10Al, may define the critical Cr and Al contents for the formation of exclusive alumina scale. External alumina scales were found on the surfaces of the ␤-NiAl and Ni–27Cr–11Al–0.5Y (wt.%) coatings, which are far above the above-mentioned line AB. The ␤-NiAl coating exhibited lower oxidation rates than the NiCrAlY coating (Fig. 5). One reason is that the NiAl coating prepared by sputtering was smoother than the NiCrAlY coating prepared by VAD (Figs. 2, 4 and 7), which causes the true surfaces of the NiCrAlY coating to be larger than the nominal surfaces. Another reason is that the Y-rich oxides formed within the alumina scale, as shown

in Fig. 8, may increase the diffusion rates through the alumina scale and therefore increase the growth rates of alumina scale [34]. One aim of addition of Y in the NiCrAlY coating is to segregate at the alumina grain boundaries and suppress grain boundary oxygen diffusion, thus lowering the scaling rate [35,36]. However, when Y rich domains of sufficient size are incorporated into the TGO, they may act as short-circuit paths for inward oxygen transport which result in an increase of the TGO growth rate [12,37,38]. No voids beneath the TGO on both the NiAl and the NiCrAlY coatings were observed. It is well known that small additions of reactive elements (REs), such as yttrium, dramatically improve the adherence of TGO to the underlying metallic substrates. The most accepted mechanism for the RE effect is that REs suppress the sulfur effect, which leads to void formation at the oxide/substrate interfaces and weakens the interface bond [39,40]. In this work, no Y was added into the NiAl coating. So, the RE effect is not applicable. Wang [18] pointed out that nanocrystallization can enhance A12 O3 scale adhesion due to: (1) reducing the thermal stress; (2) relieving growth stresses and thermal stresses by enhancing plastic deformation of the scale and the substrate, and (3) increasing the bonding force between the scale and the substrate due to the formation of oxide intrusions into the substrate grain boundaries, which act as pegs, anchoring the scale to the substrate. The ␤-NiAl coating studied in this work had fine grain sizes from 300 nm to about 1 ␮m. The thin alumina scale formed on the microcrystalline ␤-NiAl coating is adherent and free of voids at the scale/coating interface (Fig. 7f). Besides, micro-pegs along the columnar grain boundaries were also observed to be able to provide extra enhancement of the bonding strength to the scale (Fig. 10).

S. Hou et al. / Applied Surface Science 324 (2015) 1–12

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Fig. 10. Elemental mapping of ␤-NiAl coating after 50 h of isothermal oxidation at 1100 ◦ C in air.

According to the diffusional-creep theory [18], the deformation rate of a polycrystalline material in the case that grain boundary diffusion predominates can be determined by ε˙ =

B˝ıDgb d3 KT

,

(4)

where B is a numerical constant,  is the tensile stress, ˝ is the atomic volume, d is the average crystal size, Dgb is the grainboundary diffusivity, ı is the thickness of the boundary, K is the Boltzmann’s constant and T is the Kelvin temperature. According to Eq. (4), the diffusional creep rate of the scale can be strengthened by reducing the grain size of the scale. Therefore, the deformation rates of the ␤-NiAl coating with an average grain size of about 1 ␮m will be three orders of magnitude faster than those coating with grain size of about 10 ␮m, which may play an important role in relieving the oxide stresses. Furthermore, grain refinement is said to alleviate the deleterious effect of sulfur on oxidation [40,41], which will be beneficial to enhance the scale adhesion. 4.2. Inter-diffusion between the coatings and the SC superalloy substrate The SC superalloy substrate contains plenty of Co, Cr and refractory elements including W, Ta, Mo and Re, while the coatings are

simple NiAl or NiCrAlY. So, during exposure at high temperatures, inter-diffusion between the coatings and SC substrate is inevitable, because the difference of chemical compositions provide great driving force for diffusion. In this work, the as-sputtered coating contains Ni–2Al (at.%) only, and then the Al contents went down to 48 at.% after vacuum annealing at 1000 ◦ C for 1 h. The decreases of the Al contents in coating result from the inward diffusion of Al into the substrate, which leads to formation a thin IDZ (Fig. 2d). It is well known that the solidsolubility of refractory elements such as Cr, W, Mo and Re in ␥-Ni are much higher than that in ␤-NiAl and ␥ -Ni3 Al [42]. So, when the ␥-Ni phases transform into ␤-NiAl as a result of inward diffusion of Al, the refractory elements will precipitate out and aggregate into particles. The vacuum annealing of the NiCrAlY coating at the same temperature and for the same time did not lead to formation of such an obvious IDZ, because its Al contents was relatively lower, ∼20 at.%. Exposures of both the ␤-NiAl and NiCrAlY coated Re-bearing SC superalloys at 1100 ◦ C for 50 h resulted in significant growth of SRZ (Fig. 7). The formation of SRZ in Re-bearing nickel-based SC superalloys with coatings has been extensively studied [42,43]. The inward diffusion of Al from the coatings into the substrates leads to the phase transformation from ␥ to ␤ or ␥ to ␥ in the substrate. As a result, Re precipitates and forms TCP phases. Besides, without the

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S. Hou et al. / Applied Surface Science 324 (2015) 1–12

Fig. 11. Elemental mapping of NiCrAlY coating after 50 h of isothermal oxidation at 1100 ◦ C in air.

hindrance of TCP containing IDZ, the SRZ seems to form more easily [44]. It has been reported by Suzuki [45] that addition of an excessive amount of Cr into the Ni-based superalloys promotes rapid SRZ nucleation during aluminizing, probably because the precipitated Cr dissolved with Re and W to forms Cr(Re,W) phases. The TCP phase in SRZ is rich in Re but Cr deficient (Fig. 11) because of relatively low solubility of Cr in Re as referred to the Cr–Re phase diagram [46]. So there is a difference between this work and Suzuki’s. Re-rich phase precipitated at first in this work, while Crrich phase precipitated at first in Suzuki’ work. In this work, a thin Cr-rich layer at the coating/substrate interface was observed only in the NiCrAlY coated specimens. Similar results were found in a MCrAlY coated Co-based superalloy [47]. Das et al. [42] pointed out that no SRZ formed in the SC superalloys coated by Pt–Hf modified ␥–␥ coatings. In this work, however, SRZ in the NiCrAlY coated specimens was observed too. Probably, Pt plays an important role in preventing the formation of SRZ. The degradation behaviors between the NiCrAlY coated specimens and the NiAl coated ones during exposure were quite different. The ␥ –␥ phase transformation of the NiCrAlY coating is obvious (Fig. 6), and about 2/3 outer part of the NiCrAlY coating is Al-depleted (Fig. 11). In contrast, only a few ␥’ grains formed in the ␤-NiAl coating (Fig. 7f). The chemical composition range for ␥ Ni3 Al is much more narrow than that for ␤-NiAl. Furthermore, the

alumina scale formed on NiCrAlY coating was much thicker than those on the ␤-NiAl coating. The NiCrAlY coating contained 27.7 wt.% Cr and 10.3 wt.% Al while the NiAl coating contained no Cr and 48 at.% Al. The large chemical composition difference between these two coatings led to quite different inter-diffusion behavior. The inward diffusion of Cr from the NiCrAlY coating into the substrate was remarkable, which resulted in the formation of Cr(Re,W) rich phase at the original coating/substrate interface. Such a phase might act as a diffusion barrier for inter-diffusion [47]. However, the lower density of TCP needles in SRZ of the NiCrAlY coated specimens than that in the NiAl coated ones most possibly owns to that the Al contents in the former is much lower. The formation of SRZ is a major problem associated with coated Ni-based superalloys, which leads to the loss of coating and degradation of the properties of coated Ni-based superalloys, and is potentially life-limiting to turbine blades. Therefore, diffusion barrier between the coating and the superalloy substrate is needed to suppress the SRZ formation [48]. In our next article, it will be reported that an active diffusion barrier is able to completely prevent the inter-diffusion between the ␤-NiAl coating and Re-bearing SC superalloy substrate at 1100 ◦ C. The active diffusion barrier is based on an oxide thin film, which is thermodynamically less stable than alumina. So it will

S. Hou et al. / Applied Surface Science 324 (2015) 1–12

react with aluminum from both the coating and the substrate during service, and forms an alumina/metal/alumina sandwich structure which has both excellent bonding and diffusion barrier ability. The efficiency of ␣-Al2 O3 with TiN buffer layer as diffusion barrier between bond coat and bulk material of gas turbine blades was proved by Müller et al. [48]. However, ␣-Al2 O3 film can only be deposited at temperatures of 1000 ◦ C or higher because of its very high melting temperature. Besides, cracks may form in an ␣-Al2 O3 barrier layer without TiN buffer layer after annealing at 1100 ◦ C as reported by Müller et al. [48]. The crack formation might be caused by CTE mismatch and initiate at the defects, which are difficult to be avoid in the film deposition. In our previous work, the effectiveness of an arc ion platedCr2 O3 intermediate film as a diffusion barrier between NiCrAlY and ␥-TiAl was evaluated by annealing at 1000 ◦ C. The results showed that Cr2 O3 acted as an active diffusion barrier by formation of two continuous ␣-Al2 O3 layers at both the TiAl/Cr2 O3 and Cr2 O3 /NiCrAlY interfaces, suppressing the inward diffusion of Ni from NiCrAlY overlay coating to ␥-TiAl substrate effectively [49]. The active diffusion barrier may take advantage of low temperature arc ion plating at about 200 ◦ C. In addition, the thermally grown alumina layers from the Cr2 O3 layer were separated by a metallic Cr layer, which may enhance the spallation resistance. In fact, an active layer of NiCrO, instead of Cr2 O3 , will be applied to the SC superalloy/␤-NiAl coating system. This design will lead to the formation of alumina/NiCr/alumina after a shortterm high temperature treatment. NiCr is more ductile than Cr, so it may be more beneficial to the spallation resistance of the coating system with an active diffusion barrier. Much more detailed elucidations and discussions will be reported systematically in Part II. 5. Conclusions 1. Microcrystalline ␤-NiAl coating was successfully prepared on a Re-bearing SC superalloy substrate through a novel approach of magnetron sputtering followed by vacuum annealing. 2. Both the microcrystalline ␤-NiAl coating and the NiCrAlY coating formed thin and adherent ␣-Al2 O3 scales during oxidation tests at 1100 ◦ C, while the oxide scales formed on the uncoated specimens primarily consisted of spinel (Ni,Co)Al2 O4 with underlying ␣-Al2 O3 scale. A few NiTa2 O6 ridges were also observed on the uncoated specimens. 3. Inclusions of Y rich oxides were found in TGO on NiCrAlY coating, and are believed to be the main cause of much faster scaling rates of the NiCrAlY coating than those of the microcrystalline ␤-NiAl coating. 4. Both inter-diffusion zone and secondary reaction zone were observed in either NiCrAlY or NiAl coated specimens. Acknowledgements This work was supported by National key Basic Research Program of China (973 Program, No. 2012CB625100) and by National Natural Science Foundation of China (No. 51231001). References [1] H. Hiroshi, High temperature materials for gas turbines: the present and future, in: Proceedings of the International Gas Turbine Congress, November 2–7, Tokyo, 2003, pp. 1–9. [2] N.P. Padture, M. Gell, E.H. Jordan, Thermal barrier coatings for gas-turbine engine applications, Science 296 (2002) 280–284. [3] G.W. Goward, Progress in coatings for gas turbine airfoils, Surf. Coat. Technol. 108 (1998) 73–79. [4] K.K. Ma, J.M. Schoenung, Isothermal oxidation behavior of cryomilled NiCrAlY bond coat: homogeneity and growth rate of TGO, Surf. Coat. Technol. 205 (2011) 5178–5185.

11

[5] M. Seraffon, N.J. Simms, J. Sumner, J.R. Nicholls, The development of new bond coat compositions for thermal barrier coating systems operating under industrial gas turbine conditions, Surf. Coat. Technol. 206 (2011) 1529–1537. [6] T.J. Nijdam, L.P.H. Jeurgens, W.G. Sloof, Promoting exclusive alpha-Al2 O3 growth upon high-temperature oxidation of NiCrAl alloys: experiment versus model predictions, Acta Mater. 53 (2005) 1643–1653. [7] M. Daroonparvar, M.S. Hussain, M.A.M. Yajid, The role of formation of continues thermally grown oxide layer on the nanostructured NiCrAlY bond coat during thermal exposure in air, Appl. Surf. Sci. 261 (2012) 287–297. [8] J.R. Nicholls, Designing oxidation-resistant coatings, J. Miner. Met. Mater. Soc. 52 (2000) 28–35. [9] P. Niranatlumpong, C.B. Ponton, H.E. Evans, The failure of protective oxides on plasma-sprayed NiCrAlY overlay coatings, Oxid. Met. 53 (2000) 241–258. [10] Y.J. Li, Y.T. Xie, L.P. Huang, X.Y. Liu, X.B. Zheng, Effect of physical vapor deposited Al2O3 film on TGO growth in YSZ/CoNiCrAlY coatings, Ceram. Int. 38 (2012) 5113–5121. [11] V.K. Tolpygo, D.R. Clarke, K.S. Murphy, Oxidation-induced failure of EB-PVD thermal barrier coatings, Surf. Coat. Technol. 146 (2001) 124–131. [12] A.G. Evans, D.R. Mumm, J.W. Hutchinson, G.H. Meier, F.S. Pettit, Mechanisms controlling the durability of thermal barrier coatings, Prog. Mater. Sci. 46 (2001) 505–553. [13] M.W. Brumm, H.J. Grabke, The oxidation behavior of NiAl. 1. Phasetransformations in the alumina scale during oxidation of NiAl and NiAl–Cr alloys, Corros. Sci. 33 (1992) 1677. [14] H.B. Guo, T. Zhang, S.X. Wang, S.K. Gong, Effect of Dy on oxide scale adhesion of NiAl coatings at 1200 ◦ C, Corros. Sci. 53 (2011) 2228–2232. [15] D.Q. Li, H.B. Guo, H. Peng, S.K. Gong, H.B. Xu, Improved alumina scale adhesion of electron beam physical vapor deposited Dy/Hf-doped beta-NiAl coatings, Appl. Surf. Sci. 283 (2013) 513–520. [16] H. Svensson, M. Christensen, P. Knutsson, G. Wahnstrom, K. Stiller, Influence of Pt on the metal-oxide interface during high temperature oxidation of NiAl bulk materials, Corros. Sci. 51 (2009) 539–546. [17] B.A. Pint, Progress in Understanding the Reactive Element Effect Since the Whittle and Stringer Literature Review, in: P.F. Tortorelli, I.G. Wright, P.Y. Hou (Eds.), Proc. John Stringer Symposium on High Temperature Corrosion, ASM International, Materials Park, OH, 2003, pp. 9–19. [18] F.H. Wang, The effect of nanocrystallization on the selective oxidation and adhesion of Al2 O3 scales, Oxid. Met. 48 (1997) 215–224. [19] S.L. Yang, F.H. Wang, W.T. Wu, Effect of microcrystallization on the cyclic oxidation behavior of beta-NiAl intermetallics at 1000 ◦ C in air, Intermetallics 9 (2001) 741–744. [20] H.B. Guo, D. Wang, H. Peng, S.K. Gong, H.B. Xu, Effect of Sm, Gd, Yb, Sc and Nd as reactive elements on oxidation behaviour of b-NiAl at 1200 ◦ C, Corros. Sci. 78 (2014) 369–377. [21] H. Guo, D. Li, L. Zheng, S. Gong, H. Xu, Effect of co-doping of two reactive elements on alumina scale growth of ␤-NiAl at 1200 ◦ C, Corros. Sci. 88 (2014) 197–208. [22] P.Y. Hou, K.F. McCarty, Surface and interface segregation in ␤-NiAl with and without Pt addition, Scr. Mater. 54 (2006) 937–941. [23] H.Y. Lou, F.H. Wang, B.J. Xia, L.X. Zhang, High-temperature oxidation resistance of sputtered micro-grain superalloy K38G, Oxid. Met. 38 (1992) 299–307. [24] G.F. Chen, H.Y. Lou, Effect of nanocrystallization on the oxidation behavior of a Ni–8Cr–3.5Al alloy, Oxid. Met. 54 (2000) 155–162. [25] A. Rahman, R. Jayaganthan, S. Prakash, V. Chawla, R. Chandra, High temperature oxidation behavior of nanostructured Ni–Al coatings on superalloy, J. Alloy Compd. 472 (2009) 478–483. [26] D.Q. Yu, X.Y. Lu, J. Ma, S.M. Jiang, S.C. Liu, J. Gong, C. Sun, Study of oxidation behavior of the gradient NiCrAlY coating at 1000 and 1100 ◦ C, Acta Metall. Sin. 48 (2012) 759–768. [27] T.B. Massalski, H. Okamoto, P.R. Subramanian, L. Kacprzak, Binary Alloy Phase Diagrams, ASM International, Materials Park, OH, 1990, pp. 183. [28] J.A. Thornton, High-rate thick-film growth, Annu. Rev. Mater. Sci. 7 (1977) 239–260. [29] M.A. Bestor, R.L. Martens, R.A. Holler, M.L. Weaver, Influences of annealing and hafnium concentration on the microstructures of sputter deposited beta-NiAl coatings on superalloy substrates, Intermetallics 18 (2010) 2159–2168. [30] W. Zhao, Z.Q. Li, B. Gleeson, A new kinetics-based approach to quantifying the extent of metastable→stable phase transformation in thermally-grown Al2 O3 scales, Oxid. Met. 79 (2013) 361–381. [31] N. Birks, G.H. Meier, F.S. Pettit, Introduction to the High-temperature Oxidation of Metals, 2nd ed., Cambridge University Press, 2006, pp. 111–115. [32] N. Birks, G.H. Meier, F.S. Pettit, Introduction to the High-temperature Oxidation of Metals, 2nd ed., Cambridge University Press, 2006, pp. 122–124. [33] G.F. Chen, H.Y. Lou, The effect nanocrystallization on the oxide formation on Ni–11Cr–8Al alloy, Scr. Mater. 43 (2000) 119–122. [34] T. Nakagawa, I. Sakaguchi, N. Shibata, K. Matsunaga, T. Mizoguchi, T. Yamamoto, H. Haneda, Y. Ikuhara, Yttrium doping effect on oxygen grain boundary diffusion in alpha-Al2 O3 , Acta Mater. 55 (2007) 6627–6633. [35] M.C. Stasik, F.S. Pettit, G.H. Meier, A. Ashary, J.L. Smialek, Effects of reactive element additions and sulfur removal on the oxidation behavior of fecral alloys, Scr. Metall. Mater. 31 (1994) 1645–1650. [36] B.A. Pint, Experimental observations in support of the dynamic-segregation theory to explain the reactive-element effect, Oxid. Met. 45 (1996) 1–37. [37] A. Gil, D. Naumenko, R. Vassen, J. Toscano, M. Subanovic, L. Singheiser, W.J. Quadakkers, Y-rich oxide distribution in plasma sprayed MCrAlY-coatings

12

[38]

[39]

[40]

[41]

[42]

S. Hou et al. / Applied Surface Science 324 (2015) 1–12 studied by SEM with a cathodoluminescence detector and Raman spectroscopy, Surf. Coat. Technol. 204 (2009) 531–538. J. Toscano, R. Vassen, A. Gil, M. Subanovic, D. Naumenko, L. Singhelser, W. Quadakkers, Parameters affecting TGO growth and adherence on MCrAly-bond coats for TBC’s, Surf. Coat. Technol. 201 (2006) 3906–3910. L. Rivoaland, V. Maurice, P. Josso, M.-P. Bacos, P. Marcus, The effect of sulfur segregation on the adherenceof the thermally-grown oxide on NiAl – II: The oxidation behavior at 900 ◦ C of standard, desulfurized or sulfur-doped NiAl(0 0 1) single-crystals, Oxid. Met. 60 (2003) 159–178. S.L. Yang, F.H. Wang, Effect of nanocrystallization on sulfur segregation in Fe–Cr–Al alloy during oxidation 1000 ◦ C, Oxid. Met. 65 (2006) 195–205. Q. Li, X. Peng, J.Q. Zhang, G.X. Zong, F.H. Wang, Comparison of the oxidation of high-sulfur Ni–25Cr–5Al alloys in as-cast and as-sputtered states, Corros. Sci. 52 (2010) 1213–1221. D.K. Das, K.S. Murphy, S.W. Ma, T.M. Pollock, Formation of secondary reaction zones in diffusion aluminide-coated Ni-base single-crystal superalloys containing ruthenium, Metall. Mater. Trans. A 39A (2008) 1647–1657.

[43] M. Reid, M.J. Pomeroy, J.S. Robinson, Microstructural instability in coated single crystal superalloys, J. Mater. Process. Technol. 153 (2004) 660–665. [44] X.P. Tan, H.U. Hong, B.G. Choi, I.S. Kim, C.Y. Jo, T. Jin, Z.Q. Hu, Characterization of topologically close-packed phases in secondary reaction zone in a coated CMSX-4 single crystal Ni-based superalloy, J. Mater. Sci. 48 (2013) 1085–1089. [45] A.S. Suzuki, K. Kawagishi, T. Yokokawa, H. Harada, Effect of Cr on microstructural evolution of aluminized fourth generation Ni-based single crystal superalloys, Surf. Coat. Technol. 206 (2012) 2769–2773. [46] M. Venkatraman, J.P. Neumann, Bull. Alloy Phase Diagr. 8 (2) (1987) 105–107. [47] K. Zhang, Q.M. Wang, C. Sun, F.H. Wang, Preparation and oxidation behavior of NiCrAlYSi coating on a cobalt-base superalloy K40S, Corros. Sci. 50 (2008) 1707–1715. [48] J. Müller, D. Neuschütz, Efficiency of ␣-alumina as diffusion barrier between bondcoat and bulkmaterial of gas turbine blades, Vacuum 71 (2003) 247–251. [49] Y.X. Cheng, W. Wang, S.L. Zhu, L. Xin, F.H. Wang, Arc ion plated-Cr2 O3 intermediate film as a diffusion barrier between NiCrAlY and ␥-TiAl, Intermetallics 18 (2010) 736–739.