Impact of Si addition on high-temperature oxidation behavior of NiAlHf alloys

Journal of Materials Science & Technology 35 (2019) 2038–2047

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Research Article

Impact of Si addition on high-temperature oxidation behavior of NiAlHf alloys Jing Jing a , Jian He a,b,c,∗ , Hongbo Guo a,c,∗ a

School of Materials Science and Engineering, Beihang University (BUAA), Beijing 100191, China Frontier Research Institute of Innovative Science and Technology, Beihang University (BUAA), Beijing 100191, China Key Laboratory of High-temperature Structural Materials & Coatings Technology (Ministry of Industry and Information Technology), Beihang University (BUAA), Beijing 100191, China b c

a r t i c l e

i n f o

Article history: Received 17 February 2019 Received in revised form 27 March 2019 Accepted 1 April 2019 Available online 10 May 2019 Keywords: Intermetallic NiAl Rare earth element Thermal cycling Oxidation Segregation

a b s t r a c t Silicon (Si) and reactive elements (REs) like hafnium (Hf) have favorable effects on improving the high-temperature oxidation resistance. To reveal the interaction effect between them, Si with different contents was added into ␤-NiAlHf alloy in this study. Cyclic oxidation behavior at 1200 ◦ C was investigated. Results show that Si can increase the Hf-rich precipitates and suppress Hf outward diffusion, thus inhibit the RE effect leading to worse oxidation resistance. However, NiAlHf-5Si alloy exhibited a lower oxidation rate and better spallation resistance mainly because of the existence of micro-cracks and Ni2 Al3 phases which provided more rapid diffusion paths for Al3+ . © 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.

1. Introduction Thermal barrier coating systems (TBCs) play an indispensable part as an outer thermal-insulating and anti-corrosion layer on gas turbine blades [1–3]. MCrAlY (M = Ni, Co or Ni + Co) is a relatively mature alloy as the bond coat material between ceramic top coat and superalloy substrate in TBCs, which compromises bilateral different thermal expansion coefficients and provides oxidation resistance. However, as temperature exceeding 1150 ◦ C, the growth of thermally grown oxides (TGO) formed on MCrAlY accelerates and oxide spallation occurs accordingly [4,5]. Intermetallic compound ␤-NiAl is a promising alternative to MCrAlY because of its high melting temperature, low density and good isothermal oxidation resistance. Nevertheless, poor cyclic oxidation and hot corrosion resistance of ␤-NiAl limit its application as the bond coat of TBCs [6–9]. Reactive element dopants (such as Hf, Dy, Y, Ce, La) were firstly patent by Pfeil et al. in 1937 aiming at modifying chromia-forming alloys [10,11]. Stringer thought to replace “rare earth” effect with “reactive element” effect for revealing the special function of those

∗ Corresponding authors at: School of Materials Science and Engineering, Beihang University (BUAA), Beijing 100191, China. E-mail addresses: [email protected] (J. He), [email protected] (H. Guo).

slight additions [12]. Since then, tremendous investigations on the mechanisms of RE effect have been carried out. It is found that RE can not only reduce the oxide scale growth rate, but also improve the adhesion of the scale/alloy interface in many chromia-forming or alumina-forming alloys [13–16]. According to previous studies, doping moderate (0.05−0.1 at.%) REs (such as Hf, Zr, Dy) in ␤-NiAl could especially enhance scale adhesion by preventing sulfur segregation and void formation at coating/Al2 O3 scale interface. Since the REs are beneficial for oxidation-resistant alloys and coatings, there has been significant focus on practical applications of RE dopants over the past 20 years. In order to gain an insight into the RE effect in alloys and coatings further, many studies on co-doping have been implemented so far. Previous research shows that co-doping two REs can improve oxidation resistance and alumina creep retardation because of the “block effect” [17–19]. However, it is thought that adding elements such as Re, Ta, Cr, Mo to Hf doped ␤-NiAl alloy could lead to negative effect on the oxidation performance [20,21]. Besides, the interactions between REs and other impurities (such as S, C) have been discussed extensively [12,20,22,23], indicating that high Hf/C or Hf/S is beneficial for scale adhesion and Hf can weaken the negative influence by impurities effectively. In our previous work, Cr was found beneficial to the hotcorrosion resistance of RE-doped NiAl but detrimental to its cyclic oxidation behaviour [20,24]. However, when moderate amount

https://doi.org/10.1016/j.jmst.2019.04.023 1005-0302/© 2019 Published by Elsevier Ltd on behalf of The editorial office of Journal of Materials Science & Technology.

J. Jing et al. / Journal of Materials Science & Technology 35 (2019) 2038–2047 Table 1 Actual chemical compositions of as-annealed alloys identified by EPMA analysis (at.%). Alloy

Ni

Al

Si

Hf

NiAlHf NiAlHf-1Si NiAlHf-3Si NiAlHf-5Si

53.36 51.96 51.58 49.80

46.64 46.99 45.93 45.27

— 0.98 2.49 4.93

—a —a —a —a

a Hf cannot be detected effectively by EPMA due to the too low content. But according to the results by inductively coupled-plasma (ICP) analysis, the contents of Hf for the four alloys are 0.10 at.%, 0.09 at.%, 0.09 at.% and 0.11 at.%, respectively.

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of Si was added meanwhile, the adverse effect of Cr could be restrained [25,26]. The detailed mechanism is still unclear. Ma et al. prepared a NiAlHfSi coating on the single crystal superalloy IC32 by electron beam physical vapor deposition (EB-PVD) and found that the coating could effectively prevent the diffusion and oxidation of Mo from the superalloy substrate at 1150 ◦ C thus showing good oxidation resistance [27]. Other studies have been employed on the oxidation of ternary Ni-Al-Si alloys with low Al contents, which demonstrate that comparing with binary Ni–Al alloys, Si can effectively reduce the critical Al content needed for the formation of exclusive alumina scales [28,29]. In the present work, targeted to find out the effect of Si in Hf doped ␤-NiAl alloy during cyclic oxidation, alloys with different Si contents are designed, namely NiAl-0.1Hf, NiAl-0.1Hf-1Si, NiAl-0.1Hf-3Si, and NiAl-0.1Hf-5Si (in at.%). The oxidation behavior at 1200 ◦ C was investigated in detail.

2. Experimental

Fig. 1. XRD patterns of NiAlHf, NiAlHf-1Si, NiAlHf-3Si, NiAlHf-5Si alloys after annealing at 1300 ◦ C for 24 h.

Raw materials with high purity (Ni: 99.99 at.%; Al: 99.99 at.%; Si: 99.9 at.%; Hf: 99.9 at.%) were used for the preparation of Sidoped (0 at.%; 1 at.%; 3 at.%; 5 at.%) NiAlHf alloys. They were cleaned in acetone by ultrasound for 30 min and dried completely. Afterwards they were melted and cast into alloy ingots in a high vacuum (≤10−3 Pa) by vacuum non-consumable electrode arc smelting. The subsequent annealing was conducted at 1300 ◦ C in an Ar atmo˜ Pa) heat-treatment furnace for 24 h. The annealed alloy sphere (200 ingots were cut into rectangular specimens with the dimension of 12 mm × 10 mm × 3 mm. For convenience, an additional 1 mm hole was machined at the edge. Prior to oxidation tests, all the specimens were polished with SiC papers up to 800 grits and the sharp

Fig. 2. Back-scattered electron surface images of the heat-treated alloys: (a) NiAlHf; (b) NiAlHf-1Si; (c) NiAlHf-3Si; (d) NiAlHf-5Si. The insets are the high magnification images in each alloy.

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Table 2 Chemical compositions of white Hf-rich areas in Fig. 2 identified by EDS (at.%). H

Ni

Al

Si

Hf

NiAlHf NiAlHf-1Si NiAlHf-3Si NiAlHf-5Si

38.54 36.88 36.16 34.42

22.32 — — —

— 35.95 38.86 28.84

39.14 27.17 24.98 36.75

surface and cross-sectional microstructures of specimens were analysed by scanning electron microscope (SEM, Quanta 200 F). Both SEM and TEM are equipped with energy dispersive X-ray spectrum (EDS). In addition, the phase transition of preliminary alumina formed on the alloys was studied by photo-stimulated luminescence spectroscopy (LabRAM HR Evol). 3. Results and discussion

edges were round off to eliminate the edge effect. Afterwards, they were cleaned ultrasonically in acetone, dried and threaded with a short Pt wire through the hole. Cyclic oxidation tests at 1200 ◦ C were performed in a tubular furnace (GSL-1600X, China). Each cycle contained 50 min oxidation in the furnace and 10 min air-cooling outside the furnace. During oxidation, specimens were put in Al2 O3 crucibles respectively which had been calcined at 1250 ◦ C to constant weight already, and in order to make the oxygen partial pressure at the top and bottom surface identical, an M-shaped Pt wire was just placed between the specimen and the crucible base. All the weighting measurements were performed by an electronic balance (Sartorius CPA 225D, Germany) with a precision of 10−5 g, and they were divided into two steps: (1) weighing the crucible containing specimen to gain the total oxidation mass gain curve; (2) weighing the specimen alone to obtain the sample oxidation mass change curve. Each measurement was repeated three times to eliminate system and accidental errors. Three specimens for each alloy were simultaneously tested. The average values of measurements were used after the correction for the evaporation of Pt wire. The elemental constituents of alloys were obtained by electron probe micro-analyzer (EPMA, JXA-8100). The annealed alloys were observed by optical microscope (OM, Olympus BX51 M) after surface corrosion in HNO3 /HF/H2 O (volume ratio: 20:10:70) solution and by electron backscattered diffraction (EBSD, JSM-7001 F) after vibration polishing. The phases of specimens were identified through X-ray diffraction (XRD, D/max-2200) and further studied by transmission electron microscope (TEM, Technai F20). The

3.1. Alloy characterization EPMA results are shown in Table 1. It should be noted that Hf cannot be detected by EPMA due to the trace content. Inductively coupled-plasma (ICP) analysis shows that the contents of Hf are 0.10 at.%, 0.09 at.%, 0.09 at.% and 0.11 at.% for NiAlHf, NiAlHf-1Si, NiAlHf-3Si, and NiAlHf-5Si, respectively. Although there are small deviations on the contents of Si, the actual Si composition is still acceptable for further research. Fig. 1 shows the XRD results of the four as-annealed alloys. It is obvious that these alloys are mainly composed of ␤-NiAl (Space number: 221, Pm-3 m, lattice constant a ≈ 2.88 Å). It is noticed that the ␤-NiAl peaks in NiAlHf-5Si are shifted to high angle a little, indicating that the lattice distance shrinks according to the Bragg’s law. It’s quite normal for lattice reducing because of the simple substitution of Al by Si with smaller radius [30]. It is found that some small peaks (marked by ‘*’) shown in NiAlHf-5Si do not belong to ␤-NiAl, which will be discussed in TEM analysis. Fig. 2 shows the back scattered microstructures of the four asannealed alloys. The white areas represent the Hf-rich precipitates in the alloys. As can be seen, the Si-free alloy has only a few precipitates (Fig. 2(a)) [31], while the Si-doped alloys possess many Hf-rich precipitates especially along the grain boundaries (Fig. 2(b)–(d)), indicating that Si can promote the precipitation of Hf-rich phases. The specific compositions of the precipitates were analysed by EDS and the results are given in Table 2. It shows that the precipitates also contain plenty of Si. Actually, in our previous study, Si has been

Fig. 3. Metallographic morphologies of heat-treated alloys: (a) NiAlHf; (b) NiAlHf-1Si; (c) NiAlHf-3Si; (d) NiAlHf-5Si. The insets are the high magnification images in each alloy.

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Fig. 4. Oxidation kinetic curves of four alloys at 1200 ◦ C during 100 h: (a) mass gain curves, in which solid/dotted lines denote weighing results with/without crucible respectively and the inset shows mass gain of the first hour; (b) fitted lines of square mass gain (with crucible) vs. time.

proved to have a low solubility in ␤-NiAl, similarly to Hf [25]. Thus it is believed that Si competes with Hf in solubilizing in ␤-NiAl and reacts with Hf to form new compounds which prefer to precipitate. Thus the Hf-rich precipitates increase as Si is added. Optical microscope is also used for observing the microstructures of the as-annealed alloys. As shown in Fig. 3, the grain sizes of the Si-doped alloys decrease compared to the Si-free alloy. According to EBSD results, the average grain sizes are about 1 mm, 450 ␮m, 280 ␮m, and 200 ␮m, respectively for 0Si, 1Si, 3Si and 5Si doped alloys. It is because that the increasing precipitates could act as non-spontaneous nucleation sites to make grains generate easily and “drag” grain boundaries to engender finer grains. From the inset of Fig. 3(b), dotted lines can be found mainly at the grain boundaries in NiAlHf-1Si alloy, and more and finer ones can also be observed in NiAlHf-3Si alloy in the inset of Fig. 3(c). These dotted lines are supposed to be Hf-rich precipitates. Interestingly, NiAlHf-5Si displays different morphology (Fig. 3(d)) from the former alloys, in which some lamellar phases can be observed and will be discussed later, corresponding to the ‘*’ peaks shown in Fig. 1.

3.2. Cyclic oxidation behavior Fig. 4(a) gives the kinetic curves of alloys during 100 h cyclic oxidation at 1200 ◦ C. In the figure the solid and dotted lines represent the mass gains of the alloys with and without crucibles, respectively. Oxidation weight gain could be read from the solid lines, while oxide spallation could be reflected by the difference between the solid and the corresponding dotted lines. Therefore, as the solid lines show, NiAlHf-5Si has the smallest mass gain which is 0.53 mg/cm2 after 100 h oxidation (NiAlHf 0.57 mg/cm2 ), while NiAlHf-1Si and NiAlHf-3Si are oxidized much faster and the mass gain is 0.80 mg/cm2 and 0.92 mg/cm2 respectively. It should be noted that the discrepancy between the dotted and solid kinetic curves for NiAlHf-3Si is the largest, indicating that the most serious oxide spallation occurred on the surface of NiAlHf-3Si during cyclic oxidation. The inset in Fig. 4(a) shows the oxidation kinetic curves at the initial stage. Commonly, the rate of mass gain for each alloy reaches maximum during the first hour, and in which the mass gain of NiAlHf-5Si is obviously the lowest. The curves in Fig.4a follow the parabolic rule, which can be expressed by Wagner function:  w2 =kp t

(1)

Fig. 5. XRD patterns of four alloys oxidized at 1200 ◦ C for 100 h.

where w (mg/cm2 ) is the weight change per unit area; t (s) is the oxidation time; kp (mg2 /(cm4 s)) is the parabolic rate constant. According to the above equation, kp can be obtained from the fitted lines of the total mass gain squared vs. time shown in Fig. 4(b). It is clear that the kp for NiAlHf is close to that for NiAlHf-5Si, which are 8.08 × 10−7 mg2 /(cm4 s) and 6.86 × 10−7 mg2 /(cm4 s), respectively. While the kp for NiAlHf-1Si and NiAlHf-3Si is an order of magnitude higher, reaching 1.64 × 10-6 mg2 /(cm4 s) and 2.09 × 10-6 mg2 /(cm4 s), respectively. This can be partly attributed to the fast re-oxidation caused by the severe oxide spallation. Fig. 5 shows the XRD results after 100 h thermal cyclic oxidation. It turns out that four alloys maintain ␤-NiAl and the oxide scales preserve ␣-Al2 O3 till 100 h. The peaks of ␤-NiAl phase in NiAlHf-5Si are shifted to high angle a little due to the lattice contraction. The same phenomenon has been observed in Fig. 1. In addition, some small peaks of HfO2 in NiAlHf alloy can also be found in Fig. 5, which seems unreasonable because the NiAlHf alloy contains the fewest Hf-rich precipitates compared to the Si-doped ones (Fig. 2). The surface morphologies of the four alloys after 5 h, 20 h and 100 h cyclic oxidation at 1200 ◦ C are given in Fig. 6. The oxide spallation on NiAlHf-3Si is most serious that many oxide chippings appear on the surface after oxidation for 20 h and 100 h (Fig. 6(h)(i)), and the spallation on the other three alloys is negligible. Also, it can be seen that the areas with no spallation are covered by the oxide ridge network. The same ridge morphology is found on

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Fig. 6. SEM microstructure evolutions of four alloys: (a–c) NiAlHf; (d–f) NiAlHf-1Si; (g–i) NiAlHf-3Si; (j–l) NiAlHf-5Si. The insets denote higher magnification of dotted rectangular areas.

NiAlHf-5Si and the difference is that the ridges are a little denser and appear earlier. Only after 5 h oxidation, the ridges are quite visible as shown in Fig. 6(j). However, compared to NiAlHf-3Si and NiAlHf-5Si, the ridge features are not obvious on NiAlHf and NiAlHf1Si (Fig. 6(a)-(f)). The uplift amplitude of the ridges stays very low during the whole oxidation process. In addition, many heterogeneous oxides within the alloy gains can be found on the surfaces of NiAlHf, NiAlHf-1Si, NiAlHf-3Si and NiAlHf-5Si (e.g., the spots marked by the yellow circles in Fig. 6(c), (d), (g) and (j)). It is easy to know that they are all formed by the oxidation of Hf-rich precipitates but exhibit different features. As

shown in Fig. 6(c), the heterogeneous oxides on NiAlHf show crater shapes. While on NiAlHf-1Si (Fig. 6(e)) they are nodular in shape and surrounded by cracks. Compared to these, the nodular oxides on NiAlHf-3Si (Fig. 6(g)) and NiAlHf-5Si (Fig. 6(j)) are much more protuberant. And as the ridges are formed, they become the centres of the spoke-like ridges (e.g. Fig. 6(k)). There are two kinds of ridge formation mechanism. One is for the ridges along the oxide grain boundaries which are named by “intrinsic ridges” [32]. They are caused by the outward growth of oxides at the grain boundaries. As it is known, all the grainboundary diffusion is much faster than the lattice diffusion for

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Fig. 7. Back-scattered electron patterns of alloy oxide scale surface (a, c, e, g) and cross section (b, d, f, h) after 100 h cyclic oxidation at 1200 ◦ C for NiAlHf (a, b), NiAlHf-1Si (c, d), NiAlHf-3Si (e, f) and NiAlHf-5Si (g, h).

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Table 3 Chemical compositions of bright white areas in Fig. 7 identified by EDS (at.%). Area

O

Ni

Al

Si

Hf

1 2 3

57.82 53.14 57.99

1.52 4.05 2.37

36.11 30.82 15.98

— 1.21 2.72

4.55 10.78 20.93

Al3+ , leading to much faster growth of alumina at the boundaries. As a result, the oxide scale bulges to form the “intrinsic ridges”. The other is for the spoke-like ridges, named by “extrinsic ridges” [33]. With the process of the oxidation, radial micro-cracks will be formed within the grains. Then the continuous growth of ␣-Al2 O3 will heal these cracks and form spoke-like ridges. In fact, cracks similar to grain boundaries can act as short-circuit diffusion paths, and the peculiar ridge network is a vestige of these paths. After longer oxidation the two kinds of ridges can be developed into almost identical microstructure [34]. Although the NiAlHf-3Si and NiAlHf5Si have similar ridge network, they show different oxidation rates. For NiAlHf-5Si, the spallation resistance is better than NiAlHf-3Si. Fig. 7 shows the back-scattered electron images of the scale surface and cross-sectional morphologies after 100 h cyclic oxidation at 1200 ◦ C. From Fig. 7(a), ridges with small uplift amplitude can be seen on the surface. Many bright white particles are just embedded in the ridges which are identified to be Hf-rich oxides by the XRD (Fig. 5) and EDS results (Table 3). Similarly, Hf-rich oxides are visible in the corresponding cross-sectional image (Fig. 7(b)). A small amount of Hf precipitates shown in NiAlHf alloy (Fig. 2(a)) could be oxidized to form Hf-rich oxides, however, the quantity of those oxides is smaller compared to those found in Fig. 7(a) and (b). As introduced above, the ridge network is a vestige of grain boundaries, so it is reasonable that the particles are mainly formed by the oxidation of Hf which diffuses through grain boundaries to the surface. Meanwhile, it should be noted that some of the Hf-rich oxides may contribute to form the fine oxide pegs (Fig. 7(b)) at the scale/alloy interface. They can pin the oxide scale to the alloy and this is why the oxide spallation resistance of NiAlHf gets improved. On the surface of NiAlHf-1Si alloy (Fig. 7(c)), Hf-rich oxides are also found embedded in the ridges, but they are much finer. In its cross section, few Hf-rich particles can be seen (Fig. 7(d)). Instead, some coarse Hf-rich pegs are visible, since Si can react with Hf to form new compounds as mentioned above and the outward diffusion of Hf has been suppressed. There isn’t enough Hf diffusing outward, so the Hf-rich oxides in the ridges cannot nucleate and grow up continuously. Moreover, as Si contents increase, the suppression effect turns more significant. Almost no fine Hf-rich oxides can be found in the ridges of NiAlHf-3Si (Fig. 7(e)) and NiAlHf-5Si (Fig. 7(g)) scales. However, the compounds containing Si and Hf prefer to precipitate in the alloys. And these precipitates are easy to be oxidized and the products (the bright white ones in Fig. 7(e) and (g)) are namely the heterogeneous oxides shown in Fig. 6. Since these Hf-rich oxides are porous and large in size, they will act as the fast diffusion paths for oxygen to permeate, resulting in the formation of the coarse oxide pegs and the subsequent severe internal oxidation (Fig. 7(d)). In contrast to the finer ones, the coarse pegs cannot improve the scale adhesion, but instead they will cause stress concentration and act as the sources of cracks. Similar phenomenon has been found in NiAlHf-1Si, NiAlHf-3Si and NiAlHf-5Si alloys, but is partly shown in Fig. 7(d). In Fig. 7(h), the outmost oxide scale shows a jagged and curly morphology. It is certainly attributed to the dense ridge network structure. By the way, it should be noted here due to the oxide spallation the scales of NiAlHf-1Si and NiAlHf-3Si appear thinner than that of NiAlHf after 100 h oxidation. But actually the mass gains of the former two are larger (Fig. 4(a)).

3.3. Short-term oxidation behavior As described in Figs. 6 and 7, the cracks in oxide scale may have significant impact on oxidation behavior like the ridge formation, the internal oxidation and so on. To understand how the cracks are generated, short-term oxidation behavior is studied. Fig. 8 shows the luminescence spectra for alloys after 10 min, 20 min and 30 min short-term oxidation. The corresponding surface morphologies after 30 min oxidation are also given in the figure. According to the literature [35,36], ␣-Al2 O3 has a spectral doublet at nominal frequencies of 14,402 cm−1 and 14,432 cm−1 under stress-free condition, while ␪-Al2 O3 shows a spectral doublet at 14,575 cm−1 and 14,645 cm−1 . The presence of residual stress will cause shift and broadening of these doublets. At the initial high-temperature oxidation, metastable monoclinic ␪-Al2 O3 forms firstly and exhibits a needle-like or flake-like morphology. After a period of incubation time, ␪-Al2 O3 will transform to the stable hexagonal ␣-Al2 O3 which is almost granular. During the phase transition of ␪- to ␣-Al2 O3 , about 10% volume shrinkage will occur [37]. As shown in Fig. 8, after 10 min oxidation, the oxide scales of the four alloys are mainly composed of ␪-Al2 O3 . But then the difference emerges. After 20 min, only doublet for ␣-Al2 O3 can be seen for NiAlHf-3Si (Fig. 8(e)) and NiAlHf-5Si (Fig. 8(g)) alloys, indicating the transformation from ␪- to ␣-Al2 O3 has finished. While in the scales of NiAlHf (Fig. 8(a)) and NiAlHf-1Si (Fig. 8(c)) alloys, ␪-Al2 O3 is still the main phase. When the oxidation time is prolonged to 30 min, ␣-Al2 O3 without ␪-Al2 O3 can be found in the spectrum of NiAlHf-1Si, but large amount of ␪-Al2 O3 still exists in the scale of NiAlHf. The surface morphologies shown in Fig. 8b and d support the spectra results. Above all it can be concluded that Si can promote the transition of metastable ␪-Al2 O3 to steady state ␣-Al2 O3 in ␤NiAlHf alloys. Many researchers [38–40] have found that the phase transition of ␪-Al2 O3 to ␣-Al2 O3 is a diffusionless martensitic- or shear-type transformation. During the process, atoms need to rearrange by means of anion vacancies. When large reactive element ions like ˜ Å for Al3+ ), enter Hf4+ , of which the radius is about 0.8 Å [14] (0.5 into the ␪-Al2 O3 lattice, on the one hand they can distort the surrounding lattice, thus the phase transition requires more energy to overcome the distortion; on the other hand they can cause anion vacancy annihilation. Based on the two impacts, the transition of ␪to ␣-Al2 O3 is delayed. However, adding Si to NiAlHf alloys can counteract the effect of Hf. Firstly, Si will compete with Hf in solubilizing in ␤-NiAl and react with Hf to form precipitates, thus lowering the solid solubility of Hf in the alloys and reducing the effective Hf content in the lattice used for deterring the phase transition. Besides, Si4+ (about 0.4 Å) [14] is believed to act as accelerators for the scale phase transition [25,39,40], seeing that the ionic radius is small. As a result, in Si-doped alloys the transition of ␪- to ␣-Al2 O3 is much faster. About 10% volume shrinkage [37] will accompany with the transition, leading to the stress accumulation in the oxide scale and the faster the transition goes, the larger the stress is. So in 3Si and 5Si doped alloys, the much faster phase transition cause much larger stresses. In consequence, substantial radial cracks are generated as shown in Fig. 8f. Actually, many more cracks or micro-cracks are also generated in NiAlHf-5Si alloy which even appear much earlier. But the fast outward growth of Al2 O3 can promote the crack healing and leave the radial ridges as vestiges (Fig. 8(h)). As a result, NiAlHf-5Si has the lowest oxidation rate in the first hour (inset of Fig. 4(a)) and the cracks cannot extend and the spallation resistance gets improved as oxidation is prolonged. Except for the above factors, such excellent oxidation behavior of NiAlHf-5Si might also be related to the unknown second phases (Fig. 1) with lamellar microstructure (Fig. 3(d)). Therefore,

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Fig. 8. Luminescence spectra of alloys after short-term oxidation at 1200 ◦ C (a, c, e, g) and their corresponding surface morphologies after 30 min oxidation (b, d, f, h) for NiAlHf (a, b), NiAlHf-1Si (c, d), NiAlHf-3Si (e, f) and NiAlHf-5Si (g, h).

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Fig. 9. TEM analysis of as-annealed NiAlHf-5Si: (a) TEM bright-field image; (b) EDS results of circles in (a); (c) low-magnification TEM bright-field image; (d) ␤-NiAl [012] zone axis SAED pattern from (c).

TEM is performed for further study and the results are given in Fig. 9. As shown in Fig. 9(a), the second phase shows typical acicular martensite morphology. Corresponding compositions of the circle-marked zones are illustrated in Fig. 9(b). The martensite represented by the pink curve contains less Ni and more Al compared to the matrix ␤-NiAl phase (the blue curve). And obviously, large amount of C is contained in the martensite, which indicates that they are high-carbon martensites. In another region of the alloy shown in Fig. 9(c), many more martensite phases with smaller sizes and different orientations can be found. High-density phase interfaces are thus present. It is confirmed that these martensite phases are ı-Ni2 Al3 by the selected-area electron diffraction (SAED) patterns given in Fig. 9(d). And as illustrated, the hexagonal ı-Ni2 Al3 has the following orientation relationships with the cubic ␤-NiAl: ¯ ¯ 13] ¯ ı //[012]␤ and (1101) [12 ı //(100)␤ . It is well known that the high-carbon martensite possesses a twin substructure which contains plenty of twin boundaries. It is thus believed that the large number of phase interfaces as well as the twin boundaries can further promote the formation of protective alumina scale by providing fast diffusion paths. Therefore, the cracks are healed in a very short time, resulting in the reduction of oxidation rate and the improvement of spallation resistance. Moreover, since Ni2 Al3 can absorb large amount of C, the detrimental effect of the impurity on initial oxidation adhesion would be weakened. This is because that C reduction could improve the plasticity of alloys and make the residual stress more accommodated in oxide scale, hence improving the scale rumpling and reducing the scale spallation [41].

4. Conclusions In this work, NiAlHf alloys were doped with different contents of Si (0 at.%, 1 at.%, 3 at.% and 5 at.%). Their cyclic oxidation behavior at 1200 ◦ C was comparatively studied and the impact of Si was focused. Conclusions can be drawn as follows:

(1) Si preferred to form precipitates with Hf in the alloys, which increased the segregation of Hf and caused alloy grain refinement. (2) NiAlHf-1Si and NiAlHf-3Si alloys exhibited higher oxide growth rate and poor spallation resistance compared to NiAlHf alloy, while NiAlHf-5Si alloy got a little improved on these two aspects. Moreover, the oxide scales on Si-doped alloys, especially on NiAlHf-3Si and NiAlHf-5Si, showed obvious ridge network structures. (3) Si weakened the reactive element effect of Hf by suppressing its outward diffusion through gain boundaries. It accelerated the phase transition of ␪-Al2 O3 to ␣-Al2 O3 , leading to the formation of cracks or micro-cracks which in turn facilitated the generation of radial ridges and would do harm to the scale adhesion. (4) Dissimilarly, NiAlHf-5Si got improved on oxidation resistance. Because the large number of micro-cracks acted as short-circuit diffusion paths for Al3+ , resulting in fast crack healing and fast formation of protective oxide scale. The presence of Ni2 Al3 phases further promoted this process.

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