Cyclic oxidation behavior of Ni3Al-basedsuperalloy

Vacuum 169 (2019) 108938

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Cyclic oxidation behavior of Ni3Al-basedsuperalloy a,1

a,1

a,1

a

T c

a,∗∗

Jingan Li , Yuanyi Peng , Jianbo Zhang , Shan Jiang , Sheping Yin , Jian Ding , Yuting Wub, Jing Wub, Xueguang Chena, Xingchuan Xiaa,b,∗, Xin Hea, Yongchang Liub,∗∗∗ a b c

School of Material Science and Engineering, Hebei University of Technology, Tianjin, 300130, PR China School of Material Science and Engineering, Tianjin University, Tianjin, 300130, PR China Department of Precision Instruments, Tsinghua University, Beijing, 100084, PR China

A R T I C LE I N FO

A B S T R A C T

Keywords: Cyclic oxidation behavior Ni3Al-Basedsuperalloy Diffusion behavior

Cyclic oxidation behavior of a polycrystalline Ni3Al-based superalloy under as-cast and solution treated conditions was investigated in detail. SEM equipped with energy dispersive X-ray spectrometer and high-resolution transmission electron microscopy was used for microstructure observation and mechanism analysis. Results showed that oxidation process occurred along eutectic area (γ-γ′) and dual phase area (γ+γ′)interfaces firstly for as-cast alloy and oxide scales along (γ-γ′)/(γ+γ′) interfaces spalled gradually with oxidation process proceeding. As for the solution treated alloy, though oxidation process also occurred along (γ-γ′)/(γ+γ′) interfaces, no obvious spalling of oxide scales was observed under the present conditions. Due to the precipitation of Hf-containing carbides which can increase the adhesion of scales and its segregation to Al2O3 grain boundaries during cyclic oxidation to hinder outward migration of Al, leading to slower growth rate of scales. During the long-time (8–200 h) cyclic oxidation process of the solution treated alloy, O atoms continuously diffused towards the matrix along the pores in scales, resulting in thickening of Al2O3 layer. Moreover, thickening of Al2O3 layer also hindered the outward diffusion of Ni and a small amount of NiO was observed on surface of specimen after longterm cyclic oxidation, meaning the improvement of oxidation resistance.

1. Introduction Ni-based superalloy has been extensively investigated and applied in aerospace and other fields as high temperature structural materials due to its excellent specific stiffness, high yield stress, low density, high melting point and good ductility, oxidation resistance and creep resistance [1,2]. Currently, aircraft engines are designed to possess lower fuel consumption, increased thrust and efficiency, which require further improvement of the turbine inlet temperature and higher requirement of Ni-based superalloys performances [3,4]. In the past few decades, researchers focused on the investigation of superalloys. Long et al. reported the microstructure and composition of a Ni-based single crystal superalloy, and the effect of various elements on mechanical strengthening, long term stability and the oxidation resistance were classified and clarified [5]. Yang et al. indicated the effect of B element on the microstructure stability and mechanical properties of a Ni–Cr based superalloy [6]. Wu et al. studied the coarsening behavior of γ′ precipitates under different heat-treatment conditions in dual phase area

(γ'+γ) of a Ni3Al-based alloy [7]. Cao et al. revealed the microstructure homogenization and high temperature cyclic oxidation behavior of a Ni-based superalloy with high Cr content [8]. Gao et al. investigated the precipitating behavior and tensile properties of an as-cast Ni-based superalloy during heat treatment [9]. As mentioned above, current researches on Ni-based superalloys are mainly about composition optimization [5], microstructure characteristics analysis [6–9], mechanical properties [6,9], besides, some papers refer to high temperature oxidation behavior [8]. It's worth noting that high-temperature oxidation has been one of main failure modes of superalloys due to its complex and variable service environments [10], high-temperature oxidation behavior would affect element distribution, destroy microstructure stability and arise the stress concentration [11], so investigations on oxidation behavior of superalloys aim to improve its high-temperature oxidation resistance would attract more attention and become popular fields in superalloys [12]. In order to further broaden the allowed usage temperature range and adapt it for more harsh working environments, it is imminent to

∗∗

Corresponding author. Corresponding author. School of Material Science and Engineering, Hebei University of Technology, Tianjin, 300130, PR China. ∗∗∗ Corresponding author. E-mail addresses: [email protected] (J. Ding), [email protected] (X. Xia), [email protected] (Y. Liu). 1 Jingan Li, Yuanyi Peng and Jianbo Zhang contributed equally to this work and should be considered as the co-first authors. ∗

https://doi.org/10.1016/j.vacuum.2019.108938 Received 19 July 2019; Received in revised form 8 September 2019; Accepted 10 September 2019 Available online 10 September 2019 0042-207X/ © 2019 Elsevier Ltd. All rights reserved.

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characteristics of the alloy [24,25]. At present, studies on this alloy are mainly focused on development of alloys [23], microstructure evolution of γ′ phase [6,24–26], isothermal compression deformation behavior [22]. There are few reports about its high-temperature oxidation resistance behavior. It is well known that Ni3Al-based superalloy will experience oxidation when it undergoes a thermal cycling process during the service process, which is continuously from a lower temperature to higher temperature and then back to lower temperature, resulting in the reduction of strength and service life [1]. Therefore, clarifying the oxidation origin of the alloy and ascertaining the growth, expansion, spallation and secondary growth process of scales have profound significance for studying its long-time cyclic oxidation behavior. In this work, short-time and long-time oxidation behaviors of the new polycrystalline alloys under as-cast and solution treated conditions were studied, respectively. The origin of the oxides in the initial stage of cyclic oxidation, growth, expansion, spallation and secondary growth of scales were systematically analyzed. Eventually, cyclic oxidation mechanism of the alloy under both as-cast and solution treated conditions was clarified.

improve its high temperature oxidation resistance under the premise of high temperature strength. Recent developments on high-temperature oxidation resistance Ni-based alloy have been carried out. For protective scales, oxides formed by Al and Cr (Al2O3 and Cr2O3) inhibit further oxidation of the alloy [13]. Jeong et al. reported the effect of B on the oxidation behavior and scale spalling of 617 alloy at 1050 °C, and found that B segregation at the grain boundary of Cr2O3 was the reason why alloy 617B enhanced the spallation resistance [14]. Qin et al. studied the long-time (100 h) cyclic oxidation behavior of a Ni-based single-crystal superalloy with high Mo content at 1100 °C, the results indicated that continuous dense α-Al2O3 oxide scale was easily formed on the alloy with higher Al content and the volatilization of MoO3 could be prevented by reducing the diffusion of Mo [15]. Brenneman et al. reported the oxidation behavior of Ni-base superalloy GTD111 at 900 °C for 1 h–452 h, and results showed that oxidation resistance of the alloy was provided by a continuous dense Cr2O3 oxide, rather than a discontinuous Al2O3 oxide layer formed in air at 900 °C [16]. Li et al. studied the short-term oxidation behavior of IC221 M alloy at 900 °C and found that the initial oxidation mainly occurred on interdendritic region, the main oxides were NiO and a small amount of Cr2O3, ZrO2, NiCr2O4 and θ-Al2O3 [17]. Swadźba et al. studied the characterization of Al2O3 scales grown on second-generation single-crystal Ni-based superalloys during isothermal oxidation at 1050 °C, 1100 °C and 1150 °C respectively, the results showed that some θ-Al2O3 particles still exist after 100 h of oxidation at 1050 °C and 1100 °C, respectively, even in the region where the transformation to α-Al2O3was finished, the surface retained whisker morphology [18]. Interestingly, Peng et al. demonstrated that ultrafine grained Ni3Al with distributions of micronsized pores exhibited an excellent oxidation resistance at 900 °C, because the unique structure helped the alloy to grow an adherent Al2O3 scale [19]. Based on present study, it could be found that high temperature oxidation resistance of Ni-based superalloy keeps a close related to chemical composition [13], microstructure constitution [20], oxidation environment [21]. Since Ni-based superalloys are used in turbine engine combustion chambers, they are exposed to an environment without oxidation resistant coatings [4]. Therefore, it is important to study the oxidation behavior of Ni-based superalloys. However, for the superalloy, heat treatment as a conventional process would act on the above factors further impact on its oxidation resistance, which has not been discussed in-depth. In addition, various oxidation behaviors were investigated only with a fixed time and temperature, and studies of cyclic oxidation about alloys mostly focused on lower oxidation temperatures and longer time intervals to evaluate its evolution behavior [16,18,19]. Practically, the application condition of alloys is changeable, only by mastering the initial oxidation regulation can help to learn the subsequent oxidation behavior. However, to date, little attention is paid to the initial evolution behavior of cyclic oxidation and the oxidation characteristics of different micro areas, which are not conducive to optimize and regulate high temperature oxidation resistance of superalloy. Therefore, it is necessary to investigate the effect of heat treatment on oxidation behavior and the initial oxidation behavior of superalloy. Recently, a new polycrystalline Ni3Al-based superalloy (possessing excellent comprehensive performances) on the basis of IC396 series alloy was developed by adjusting the elements of W, Mo, Ti and Hf [22,23]. The alloy consists of dual phase area (γ+γ′) with a volume fraction of about 80% and eutectic area (γ-γ′) of 20% around. Moreover, (γ+γ′)/(γ-γ′) interface areas (in a non-equilibrium state due to the micro-segregation during casting process) represent typical structural

2. Materials and procedures 2.1. Specimen preparation Polycrystalline Ni3Al-based alloy was produced by vacuum induction melting (VIM) and electro slag remitting (ESR). To improve the metallurgical quality and keep the chemical composition stable, the vacuum pressure was chosen as 1 × 10−2 Pa and the melting temperature was kept at 1600 °C. Based on the aviation industry standard of HB 5220-2008 [27], carbon and sulfur contents were about 1 × 10−3 and 3 × 10−5 through high frequency induction combustion-infrared absorption method, oxygen and nitrogen contents were about 5 × 10−6 and 7 × 10−6 respectively through impulse heating-infrared and thermal conductivity method, hydrogen content was about 1 × 10−5 through impulse heating thermal conductivity method, phosphorus content was about 1 × 10−4 through n-butyl alcohol-chloroform extraction photometric method. Meanwhile, contents of other impurity elements were also determined according to HB 5220-2008. In order to eliminate casting stress, the alloy was homogenized at 850 °C for 2 h (named as the as-cast alloy unless otherwise noted). The detailed composition of the alloy used in this research is listed in Table 1. Specimens were cut into dimension of 8 mm in diameter and 8 mm thick from the same position of the as-cast alloy to ensure chemical composition consistency, and ground with sand disk up to 1500 grit. Where after all specimens were ultrasonically cleaned in alcohol for 10min and dried by cold flowing air. It is noteworthy that no oxide on the surface of specimens was observed before the cyclic oxidation experiment. Each specimen was put in a corundum crucible with the dimensions of 20 mm × 20 mm × 15 mm (length × width × height) for cyclic oxidation tests. It should be noted that the crucibles were preheated for more than 100 h at 1100 °C to confirm the weight is stable in the test environment. High temperature box resistance furnace with a maximum service temperature of 1300 °C and accuracy of ± 2 °C were used for cyclic oxidation tests. The polished as-cast specimens were taken for solution treatment at 1200 °C for 8 h and polished the specimen surface to remove the small amount of scales. The as-cast and solution treated specimens were weighed using electronic balance with accuracy of ±

Table 1 Composition of JG4246A alloy.(wt.%). Cr

Al

Ti

W

Mo

Hf

C

Si

Fe

Mn

B

Ni

7.4–8.2

7.6–8.5

0.6–1.2

1.5–2.5

3.5–5.5

0.3–0.9

0.06–0.2

< 0.5

<2

< 0.5

< 0.05

Bal.

2

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oxidation are plotted in Fig. 2. Fig. 2a shows the weight change of specimens after cyclic oxidation at 1100 °C for 200 h. It shows that the weight of specimens gradually increases as the oxidation progresses. After oxidized for 8 h, the weight of the as-cast Ni3Al alloy increases by 0.6967 mg/cm2, which is obviously higher than the mass gains of the solution treated alloy (0.4976 mg/cm2). It is clear that in the initial stage of the oxidation, the solution treated alloy exhibits better oxidation resistance than the as-cast alloy. With the oxidation process proceeding, weight change of the as-cast alloys declines firstly with cycle oxidation time going on, and then shows a tendency of increase-decrease-increase (Fig. 2a). Weight change of the solution treated alloy shows same fluctuation trend, but it should be point out that there are three mass gain platforms (as the blue arrows indicated) during cyclic oxidation for 200 h compared with the as-cast alloy, indicating that the growth and spallation behaviors of the oxide reaches an equilibrium state. This means that although the weight change has a tendency to fluctuate up and down the solution treated alloy is more stable than the as-cast alloy under long-time cyclic oxidation. As shown in Fig. 2b, weight change of the as-cast and solution treated alloys at 1100 °C for 8 h are greatly different from the oxidation kinetics curve of long-time cyclic oxidation. This means that the as-cast and solution treated alloys have different oxidation behaviors in the initial stage of oxidation, which is of great significance for the study of long-time cyclic oxidation mechanism. Therefore, results of cyclic oxidation time for 8 h are extracted and plotted in Fig. 2b to analyze the short-time oxidation behavior of the as-cast and solution treated alloys. For the as-cast alloy, mass gain occurred with oxidation time of 10min. Obvious mass loss occurred with oxidation time of 1 h due to severe spallation of scales, which could also be seen by naked eye. As the cyclic oxidation proceeded, surface of the alloy was continuously oxidized and the oxide grew up gradually, which leads to mass gain. During the subsequent cyclic oxidation process, the weight change curve shows slight fluctuations around the equilibrium position. Compared to the ascast alloys, solution treated alloys maintain a relatively stable trend in the initial stage of the cyclic oxidation. There is no significant weight change within the oxidation time of 8 h due to the small amount of oxides growth and spallation. Notably, the mass change after cyclic oxidation for 10min is slow, and mass loss don't begin until the cyclic oxidation time exceeds 0.5 h, which is completely different from the initial oxidation behavior of the as-cast alloy. Overall, the solution treated alloy has a better oxidation resistance, especially in inhibition of transient oxidation behavior in the initial stage.

0.0001 g as the initial weight of the alloy. After that as-cast and solution treated specimens were put into the homogeneous temperature zone of the box resistance furnace with fixed temperature of 1100 °C respectively. Specimens used for long-time cyclic oxidation were taken out every 8 h then cooled to room temperature in air, and this process was repeated for 25 times, meaning 200 h were used for long-time cyclic oxidation and specimens after cyclic oxidation for 40 h and 80 h were taken out for microstructure observation. Specimens used for short-time cyclic oxidation were taken out every 1 h, similarly, cooled to room temperature in air, and this process was repeated for 8 times, meaning 8 h were used for short-time cyclic oxidation and specimens after cyclic oxidation for 0.5 h、1 h、2 h、4 h、6 h and 8 h were taken out for microstructure observation. Each specimen for microstructure observation was weighed by electronic balance. In addition, the as-cast and solution treated specimens were still oxidized in air at 1100 °C for 10min, 20min for initial cycle oxidation, and air-cooled to room temperature to investigate the initial oxidation behavior. 2.2. Characterization X-ray diffraction (XRD, Bruker D8 Advance), which used a diffractometer coupled with Cu Kα as radiation, and an applied current and voltage of 40 mA and 40 kV, was carried out on the surfaces of the oxidized specimens to analyze the composition of scales. The scan angle is between 20° and 80°. The surface and cross-sections of the specimens were observed by field-emission scanning electron microscopy (FESEM, Nova Nano SEM450, the vacuum degree is 6 × 10−4 Pa). Chemical composition was determined using energy-dispersive spectroscope (EDS) equipped in the scanning electron microscope. After characterizing all specimens scales, the surface of oxidized specimens for 10min and 20min were polished, typical metallographic preparation process and Nimonic etching solution were applied for microstructure observation. Microstructure of (γ-γ′)/(γ+γ′) interfaces were characterized by SEM equipped with EDS. 3. Results 3.1. Microstructure of as-cast and solution treated Fig. 1 shows the typical microstructures of as-cast and solution treated alloys, the alloys are mainly composed of about 20 vol % eutectic area (γ-γ′) and 80 vol % dual phase area (γ+γ′). The detailed description about microstructure of the as-cast alloy has been introduced in our previous study [24,25]. Compared with the as-cast alloy, the solution treated alloy presents a different microstructure (Fig. 1b), the morphology of eutectic area (γ-γ′) is relatively uniform and the contour which has sharp corners in the as-cast alloy becomes smoother.

3.3. Short-time cyclic oxidation 3.3.1. Initial oxidation stage In order to further investigate the initial oxidation behavior of the as-cast and solution treated alloys, samples with oxidation time of 10min and 20min were selected for detailed analysis. XRD detection results are shown in Fig. 3. It can be seen that various types of oxides formed during the oxidation tests at 1100 °C. Analysis on the as-cast alloy pattern reveals that Ni3Al phase is the dominant phase at the

3.2. Oxidation kinetics The oxidation kinetics curves for long-time and short-time cyclic

Fig. 1. (a) Backscatter SEM images of microstructure of as-cast alloy, (b) secondary SEM images of microstructure of solution treated alloy. 3

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Fig. 2. Cyclic oxidation kinetic curves: (a) Long-term, (b) Short-term.

and dual phase area (γ+γ′) oxides are shown in Fig. 4f and g, respectively. It can be found that oxides of the dual phase area (γ+γ′) are mainly the mixture of NiO and Al2O3. It is worth noting that the content of Al2O3 in the eutectic area (γ-γ′) is higher, which is mainly due to the micro-segregation during the casting process resulting in a higher Al content in this area [5]. Fig. 4c shows surface morphology of the as-cast alloy cyclic oxidized at 1100 °C for 20min. As the oxidation progresses, the oxide (mixture of NiO and Al2O3) begins to expand along (γ-γ′)/ (γ+γ′) interfaces to both sides, and partially forms NiAl2O4 spinel [13]. Besides, a small amount of white granular and elongated oxide distributes in local areas of the (γ-γ′)/(γ+γ′) interface, EDS analysis (Fig. 4h) shows that the oxides are rich in Cr (mixture of Cr2O3 and NiCr2O4). The region including eutectic area (γ-γ′), dual phase area (γ+γ′) and (γ-γ′)/(γ+γ′) interface are enlarged in Fig. 4d. After the ascast alloy is cyclic oxidized at 1100 °C for 20min, NiO in (γ-γ′)/(γ+γ′) interface begins to peel off. With cyclic oxidation time reaching to 20min, oxides in eutectic area (γ-γ′) and dual phase (γ+γ′) area significantly increase. Fig. 5a presents surface morphology of the solution treated alloy oxidized at 1100 °C for 10 min. Compared with the as-cast alloys, small quantity of oxides exist on the surface during the initial stage of

initial state (Fig. 3a). For the as-cast alloys, oxides on the surface of the alloy with cyclic oxidation time of 10min are composed of NiO, Al2O3 and spinel NiAl2O4. The diffraction peaks of Cr2O3 and NiCr2O4 appear after cyclic oxidation at 1100C for 20min. For the solution treated alloys, after initial stage of cyclic oxidation (10min, 20min), little NiO is found when oxidized for 10min, while Al2O3 oxides account for the majority (Fig. 3b). Increasing oxidation time to 20min, diffraction peaks intensity of Al2O3 increases accompanied by the appearance of NiAl2O4 phase. Fig. 4 and Fig. 5 illustrate the surface morphology of the as-cast and solution treated alloys during the initial cyclic oxidation. For the as-cast alloy significant oxidation occurs along (γ-γ′)/(γ+γ′) interfaces area (Fig. 4a) between eutectic area (γ-γ′) and dual phase area (γ+γ′). Under higher magnification of (γ-γ′)/(γ+γ′) interface, the oxide is composed of a large amount of gray smooth nano-particles (Fig. 4b). Fig. 4e shows the EDS results on the (γ-γ′)/(γ+γ′) interface oxide after oxidation at 1100 °C for 10min indicate the oxide is mainly NiO. In addition, eutectic area (γ-γ′) and dual phase area (γ+γ′) are also slightly oxidized, where the size of oxide particles is obviously smaller than that of NiO particles, besides, it could be observed the interface of particles is sharp and appears dark gray. EDS detection results on the eutectic area (γ-γ′)

Fig. 3. Phase composition of Ni3Al-based superalloy after cyclic oxidation at 1100 °C for 10 min, 20 min (a) as-cast alloy, (b) solution treated alloy. 4

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Fig. 4. Secondary SEM images of surface morphology and EDS analysis of as-cast alloy after cyclic oxidation at 1100 °C for 10 min and 20 min (a) 10 min, (b) magnified image of the (γ-γ′)/(γ+γ′) interface after 10 min, (c) 20 min, (d) magnified image of the (γ-γ′)/(γ+γ′) interface after for 20 min, (e) EDS analysis of (γ-γ′)/ (γ+γ′) interface area oxides, (f) EDS analysis of eutectic area (γ-γ′) oxide, (g) EDS analysis of dual phase (γ+γ′) oxides, (h) EDS analysis of white phase in the (γ-γ′)/ (γ+γ′) interface area.

occurs. Surface morphology of the solution treated alloy which was cyclic oxidized at 1100 °C for 20min is shown in Fig. 5c. With cyclic oxidation time increasing to 20min, part black discontinuous point oxide (mixture of NiO and Al2O3in blue box) in (γ-γ′)/(γ+γ′) interfaces spalls gradually because of different thermal expansion coefficient between eutectic area (γ-γ′) and dual phase area (γ+γ′). In addition, the oxides distributing in eutectic area (γ-γ′) and dual phase area (γ+γ′) began to increase, but the oxide is significantly less than that of the ascast alloys which are oxidized at 1100 °C for 10min.

oxidation and surface of the specimen are relatively smooth. As it could be seen from the enlarged view of Fig. 5b, oxidation preferentially occurs along the (γ-γ′)/(γ+γ′) interface and forms black discontinuous point oxide as indicated by the red arrows. It is obvious that these point oxides evenly distributes along (γ-γ′)/(γ+γ′) interfaces and EDS results show that the point oxides are rich in O, Ni and Al elements, indicating they are mixture of NiO and Al2O3 (Fig. 5e) [21]. Besides, a small amount of irregular bright white phase distributes at the cusp of partial (γ-γ′)/(γ+γ′) interfaces. EDS results (Fig. 5f) show that mass percentage of Hf element in the white phase is about 60.86%, meaning the irregular bright white phase is mainly Hf-containing carbides (Fig. 5b). Nevertheless, the entire morphology of surface does not change markedly compared to initial microstructure, suggesting slight oxidation

3.3.2. Short-time oxidation stage In order to prove the effect of initial oxidation behavior on subsequent cyclic oxidation, the as-cast and solution treated specimens 5

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Fig. 5. Secondary SEM images of surface morphology and EDS analysis of solution treated alloy after cyclic oxidation at 1100 °C for 10 min and 20 min (a) 10 min, (b) magnified image of the (γ-γ′)/(γ+γ′) interface after oxidation for 10 min, (c) 20 min, (d) magnified image of the (γ-γ′)/(γ+γ′) interface after oxidation for 20 min, (e) EDS analysis of black discontinuous point oxides, (f) EDS analysis of bright white phases.

containing a small amount of large spherical NiO oxide according to the EDS results. NiO particles are nanometer particles (amplified in inset image of Fig. 7b). With the cycle oxidation time increasing, spherical NiO particles gradually grow, but some of them spall from the surface of specimen (as indexed by yellow dotted box in Fig. 7e and f, and the dotted box is the position NiO falls off). Under higher magnification of rough area (inset image in Fig. 7c), Al2O3 densely distributes on the surface of specimen, and a small amount of sharp complex oxide consisting of NiO and NiAl2O4 distributes in local area, which has finer and smaller grains [13]. It can be clearly seen that as the oxidation proceeds, some small oxide particles grow on the dense Al2O3, meaning that former oxides falling off and new oxides gradually growing during the process of cyclic oxidation. For solution treated alloy (Fig. 7g-l), after cyclic oxidation for 0.5 h, irregular oxides formed on the surface of the alloy, and micro-cracks appear at the top of local oxide. The inset image in Fig. 7g shows the microstructure of oxide on partially smooth surface. A large amount of white flower-like oxide evenly distributes on the surface of the alloy. With cyclic oxidation reaching to 1 h, spallation behavior occurs around the irregular oxide and irregular oxides become spherical NiO. The inset image in Fig. 7g shows the morphology of the oxides on the smooth region around the spherical oxide, which is mainly composed of representative dense Al2O3 oxide. When cyclic oxidized for 2 h, the area around spherical NiO re-forms new oxide and is bonded to the spherical NiO. When the cycle oxidation time is 4 h, the spherical NiO gradually grows up. There are some cracks among the larger spherical NiO particles (Fig. 7j), resulting in the bonding force between particles and the matrix weakened, and the scaly exfoliation

with cyclic oxidation time of 8 h were studied. XRD detection results of both the as-cast and solution treated alloys cyclic oxidized at 1100 °C for 8 h are shown in Fig. 6. When oxidation reaches 0.5 h, it is indicated that diffraction peaks of Cr2O3 and NiCr2O4 completely disappeared, only the diffraction peaks of the matrix, NiO, Al2O3 and NiAl2O4 can be observed. Increasing cyclic oxidation time to 2 h, the content of NiO and Al2O3 in the as-cast alloys increases rapidly, moreover, peaks of Al2O3 are stronger than that of NiO (Fig. 6a) indicating the oxides increases and Al2O3 accounts for the majority of all oxides. When the alloy was cyclically oxidized for 4 h, diffraction peaks of NiO and Al2O3 decreased gradually and the scale peeled off during this process, and the same trend appears on NiAl2O4. During short-term cyclic oxidation process (8 h), the peak of the oxide is detected to rise and then decrease for several times, which reflects the continuous growth and spallation behaviors of oxides. Compared with the as-cast alloy, each diffraction peaks intensity of the solution treated alloy after cyclic oxidation is significantly lower than that of as-cast alloy and peaks of only NiO, Al2O3 and NiAl2O4 are detected (Fig. 6b). As the cyclic oxidation proceeds, change of the oxides diffraction peak intensity is consistent with that of the as-cast alloy. It should be noted that the peak of NiAl2O4 disappears after 6 h of cyclic oxidation and a sharper peak of Al2O3 is detected. Increasing cyclic oxidization time to 8 h, the intensity peak of NiO decreases. Obviously, the peak of Al2O3 is the main part so that it could be regarded the main oxide on the surface. Fig. 7 shows surface morphology of the as-cast and solution treated alloys cyclic oxidized at 1100 °C for 8 h. For the as-cast alloy (Fig. 7a–f), obvious oxide appears on the surface after cyclic oxidation for 0.5 h, 6

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Fig. 6. Phase composition of the alloys cyclic oxidation at 1100 °C for 0.5 h, 1 h, 2 h, 4 h, 6 h, 8 h: (a) as-cast, (b) solution treated.

the spallation of outer layer NiO and the continuous growth of inner layer of Al2O3, which is consistent with the results of surface scanning results (Fig. 7).

phenomenon appears as shown by the yellow arrow. When the cyclic oxidation time reaches 8 h, surface of the solution treated alloy is substantially covered by denser and flat Al2O3 with higher oxidation resistance, occasionally small amount of spherical NiO particles present (Fig. 7l). In order to further illustrate the oxidation resistance of Ni3Albasedalloy after different heat treatment, Fig. 8 shows a series of secondary electron images taken from cross-sections of multilayered scales for different cyclic oxidation times. Two individual layers can be clearly distinguished after a certain cyclic oxidation time. For the as-cast alloy, when the alloy is cyclically oxidized for 4 h, the oxide layer is mainly composed of inner oxide layer and outer oxide layer (Fig. 8a). Surface scanning of selected area in yellow box reveals that main component of inner oxide layer was Al2O3 (Fig. 8d). Due to spallation of partial oxide, surface scanning cannot accurately characterize the composition of the outer oxide layer. Combined with the EDS results in Fig. 8b, the outer oxide layer is mainly NiO. As the cyclic oxidation progresses, the thickness of inner oxide layer and outer oxide layer gradually increase. After cyclic oxidation of 8 h, most of outer oxide layers cracks, and only a small amount of NiO oxide exists in local area. For the solution treated alloy, Fig. 8e–g shows light different morphologies compared with the as-cast alloy, though the thickness of oxide layer decreases and degree of fold decreases is observed. The surface scanning analysis about the cross section of cyclic oxidation for 4 h shows that the inner oxide layer of the alloy after solution treatment is still Al2O3 (Fig. 8h). It is obvious that NiO layer significantly decreased, and some NiO particles exist on the Al2O3 layer as indicated by the red arrow in Fig. 8g. After cyclic oxidation for 8 h, the surface becomes smoother, which may be due to

3.4. Long-time cyclic oxidation As described above, the solution treated alloy has superior oxidation resistance than the as-cast alloy under short-time cyclic oxidation conditions. In this section, long-time cyclic oxidation behavior of the solution treated alloys was studied in detail to determine the long-term serviceability of the alloy. XRD detection results of the solution treated alloy during long-time cyclic oxidation (40 h, 80 h, 120 h, 160 h, 200 h) at 1100 °C are shown in Fig. 9. It can be seen that diffraction peaks of the solution treated alloy after long-time cyclic oxidation are mainly composed of the matrix (Ni3Al) and Al2O3, and no diffraction peaks of other oxides were observed under the present conditions. When cyclic oxidation time is 40 h, diffraction peaks of Al2O3are lower. When cyclic oxidation time is 80 h the diffraction peak intensity increases, suggesting the content of protective Al2O3 increases. When the alloy is cyclically oxidized at 1100 °C for 120 h, the diffraction peaks of Al2O3 slightly decrease, indicating weakening of the bond between Al2O3 and matrix during this process leads to the spallation of the oxide. When the alloy is cyclically oxidized for 200 h, the diffraction peak intensity of Al2O3 is the highest, indicating that scale continues to grow and a new oxidation begins. Fig. 10 shows surface morphology of the solution treated alloys after long-time cyclic oxidation (40 h, 80 h, 120 h, 160 h, 200 h). It is clear that after 40 h of cyclic oxidation, bright gray scale-like oxides appear 7

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Fig. 7. Secondary SEM images of surface morphologies of the alloys: (a) as-cast-0.5 h; (b) as-cast-1h; (c) as-cast-2h; (d) as-cast-4h; (e) as-cast-6h; (f) as-cast-8h; (g) solution treated-0.5 h; (h) solution treated -1 h; (i) solution treated-2h; (j) solution treated-4h; (k) solution treated-6h; (l) solution treated-8h.

continuously decreases, which is consistent with the weight change curve of long-time cyclic oxidation of the alloy.

on the surface (Fig. 10a), and EDS results show the oxide is Al2O3, which is consistent with XRD detection results. The surface of alloy is relatively flat and a small amount of pores exists at local areas. When cyclic oxidation time reached to 80 h, the oxide scale (Fig. 10b) on the surface grows closely, the oxide grows continuously near the pores, resulting in a greatly reduced porosity. With the progressing of cyclic oxidation (Fig. 10c), Al2O3 oxide warps and spalls at various places, leading to ridge-like interconnected morphology. When the alloy is cyclically oxidized for 160 h (Fig. 10d), peeling off behavior in large area of oxide scale occurs as shown by the white dotted line. Fine oxides of the inner layer exposed because of peeling off of oxide area, and EDS results show that the oxide is Al2O3. A thicker Al2O3 protective layer formed on the surface of the alloy after a long period of cyclic oxidation. In addition, bright white flower-like Al2O3 oxide scale concentrates on the right side of Fig. 10d. As marked by the yellow arrow, cracks appear in the vicinity of the flower-like Al2O3, causing warping of oxide scales at various places. Fig. 10e shows the microscopic morphology of the oxide scale after cyclically oxidized at 1100 °C for 200 h. With the cyclic oxidation time increasing, oxides in the spallation region of alloy surface grow again, so that the porosity of the oxide layer

4. Discussion 4.1. Initial oxidation of (γ-γ′)/(γ+γ′) interface As described above, oxidation behavior between the as-cast and solution treated specimens are significantly different during the initial stage of cyclic oxidation. The as-cast specimens for cyclic oxidation of 10min and 20min were investigated in detail (Figs. 3–5). XRD results show that the oxides on the surface of as-cast alloys are composed of NiO, Al2O3 and spinel NiAl2O4. It is known that NiAl2O4 is formed through the reaction of (1) [ [28,30–33]]. Al2O3+Ni+1/2O2→NiAl2O4 or Al2O3 + NiO→ NiAl2O4

(1)

In other way, NiAl2O4 transforms to Al2O3 by reaction of (2) [28,34]]. 3NiAl2O4 +2Al → 4Al2O3 + 3Ni 8

[

(2)

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Fig. 8. Secondary SEM images of cross section microstructures of the alloys after short-term cyclic oxidation: (a) as-cast-4h, (b) as-cast-6h, (c) as-cast-8h, (d) surface scanning results of as-cast-4h, (e) solution treated-4h, (f) solution treated-6h, (g) solution treated-8h, (h) surface scanning results of solution treated-4h.

are stronger, evidencing Cr-containing oxides are unstable at this temperature, and NiCr2O4 is thought to be formed by a reaction of (3) [ [16,35–40]]. Cr2O3+Ni+1/2O2→NiCr2O4 or Cr2O3+NiO→NiCr2O4

(3)

Existence of NiCr2O4 means low stability of Cr2O3 or high activity of oxygen in the scale, leading to a poor oxidation resistance [49]. It can be seen from Fig. 3b that when the solution treated alloy was oxidized at 1100 °C for 10 min, little NiO is observed and Al2O3 is the main content of the oxide according to the relative peak strength. Therefore, oxidation resistance of solution treated alloys is improved due to the presence of a large amount of protective Al2O3 compared to the as-cast alloy. With oxidation time prolonging to 20min, the diffraction peaks intensity of Al2O3 increased and NiAl2O4 formed. In addition, Cr2O3 appeared during the cyclic oxidation at 1100 °C for 20min (Fig. 3b) and the content of Cr2O3 was lower than that of the as-cast alloy. The absence of NiCr2O4 in the initial stage of oxidation contributes to better oxidation resistance [29]. Based on previous research, it has been proved that due to the casting process of the Ni3Al-based superalloy, Al element mainly segregated in the eutectic area (γ-γ′) and its content along the interface area is relatively lower [25]. Higher internal energy and unstable state around the (γ-γ′)/(γ+γ′) interfaces will cause crystal defects such as vacancies and dislocations during high temperature (1100 °C) oxidation process [24,25,41,42], which can provide channels for fast diffusion of atoms [43]. Therefore, Ni in the matrix will preferentially diffuse into the interface area and the diffusion rates of atoms in the (γ-γ′)/(γ+γ′) interfaces are faster than that in other areas. Meanwhile, due to the lower Al content in the interfacial area, formation of Al2O3 on this area is rare. As shown in Fig. 4b, the (γ-γ′)/(γ+γ′) interfaces is covered by

Fig. 9. Phase composition of the alloys solution treated at 1100 °C for 40 h, 80 h, 120 h, 160 h, 200 h.

With the increasing of oxidation time (20min), the peak intensity of Al2O3 increases significantly, indicating that the oxide film becomes thicker. In addition, the relative intensities of NiO peaks increase slowly with oxidation time increasing compared to that of Al2O3 peak as indicated in Fig. 3a, suggesting that part of NiO peels from the surface during rapid cooling, resulting in its content decrease. Moreover, Cr2O3 and NiCr2O4 peaks appear after the cyclic oxidation at 1100 °C for 20min and the peaks of Cr2O3 are very weak and the peaks of NiCr2O4 9

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Fig. 10. Secondary SEM images of surface morphologies of solution treated alloy after long-term cyclic oxidation: (a) 40 h, (b) 80 h, (c) 120 h, (d) 160 h, (e) 200 h

NiO to start forming an outer oxide layer. From the above, NiO first formed on the surface of the specimen in the initial stage of cyclic oxidation, which shows good agreement with the other researchers results [15,19,44]. Higher content of Al element in the eutectic area (γγ′) makes Al2O3 oxide form during high temperature oxidation [18]. When the as-cast alloy is cyclic oxidized at 1100 °C for 20 min, NiO existed along the (γ-γ′)/(γ+γ′) interfaces (Fig. 4d) begins to peel off. Section 4.2 describes the specific spalling behavior of scale in detail. In addition, Cr also diffuses from the dual phase area (γ+γ′) to the (γ-γ′)/ (γ+γ′) interfaces due to its concentration gradient during the oxidation process [49]. The dislocations and microcracks at the interfaces provide diffusion channel for Cr element, so that Cr concentrates in this region to form a Cr-rich phase [43]. Further, Cr2O3 act as protective barrier when the temperature is up to about 1600 °F (871 °C), so the oxidation resistance of the as-cast alloy decreases as the Cr-rich phase is oxidized to be Cr2O3 at the (γ-γ′)/(γ+γ′) interface (Fig. 4d). Cr2O3would be unstable above this temperature but Al2O3 dominates the protection [45]. Therefore, volatilization or decomposition of Cr2O3 oxide at 1100 °C leads to a significant increase in oxidation rate [14,46]. After solution treatment at 1200 °C for 8 h, Cr and Al elements uniformly distribute in the alloy to eliminate micro-segregation relatively [47]. The solution treated alloy is oxidized for 10 min, NiO forms around the (γ-γ′)/(γ+γ′) interfaces as shown in Fig. 5b. Al2O3 oxides also formed during the oxidation process due to the increase of Al content around the (γ-γ′)/(γ+γ′) interface areas after solid solution treatment. It is reported that the oxidation rate of Al is much faster than that of Ni [48] and the affinity between Al and O is stronger than that of Ni [50]. As the oxidation proceeding, Al atom in the matrix diffuses continuously to the surface of the alloy, resulting in rapid growth of Al2O3 at (γ-γ′)/(γ+γ′) interface (Fig. 5b). Therefore, oxides in the (γγ′)/(γ+γ′) interface region are mainly composed of Al2O3 and a small amount of NiO. In addition, solution treated alloy has a high interface energy, especially on the cusp of the interface, and the nearby atoms arrange chaotically [36]. During cooling process, exsolution of solute atoms (Hf) with larger radius first occurred and preferentially diffuse toward the cusp of the interface to reduce the total energy of the system [49]. It is known that the growth of Al2O3 scale is mainly affected by grain boundary diffusion [51]. In this process, Hf cations tend to segregate on Al2O3 grain boundaries at the (γ-γ′)/(γ+γ′) interfaces during cyclic oxidation [52]. Wang et al. observed Hf-rich grain boundaries on the surface of Hf-modified aluminide coatings [53]. The enrichment of

Hf at the grain boundary blocks the outward migration of Al, leading to slower growth rate of scales. In addition, the precipitation of Hf-containing carbides can increase the adhesion force of scales [54], which is one of the key factors to improve oxidation resistance. As the oxidation proceeding, partial oxides begin to spall along the (γ-γ′)/(γ+γ′) interface due to the difference of thermal expansion coefficient (Fig. 5d) [44].

4.2. Growth and spalling behavior of scale during cyclic oxidation When the as-cast alloy is cyclically oxidized at 1100 °C for 20 min (Fig. 4d), NiO on the (γ-γ′)/(γ+γ′) interfaces begins to spall. It is reported that for spontaneous spallation of an oxide from a matrix, a preexisting crack or separation at the oxide layer/matrix interface is required [55,56]. Cracks could generate by rumpling of the interface [56,57]. Because the oxide layer at room temperature is much stiffer than the matrix, during the cooling process, the thermal expansion and contraction of the matrix cause the oxide layer fail to follow the deformation of the matrix, resulting in cracks at the interface [55]. In addition, according to Park et al., cracking mainly resulted from the large difference in the thermal expansion coefficients [58,59]. NiO at the (γ-γ′)/(γ+γ′) interfaces has even larger thermal expansion coefficient of 17.1 × 10−6 K−1 [60], thus spallation can occur easily if NiO is connected with other oxides (such as Al2O3, NiAl2O4). With the oxidation progressing, the oxide on specimen surface grows continuously and the scales gradually thicken. It can be inferred from the aforementioned results that the oxide film of the present polycrystalline Ni3Al-based alloy exhibits a two-layer structure, which is characterized by a NiO dominant external layer and Al2O3 inner layer. The two-layer structure is closely connected at short-time oxidation, while longer time of oxidation gives rise to local peeling off of external layer, leading to the roughening of surface. Fig. 11 presents the overall description of transient oxidation. The formation as proposed elsewhere [61,62] of oxide is due to the outward diffusion of Al along the “short-circuit” twin boundaries in the oxide. It has been reported [63] that oxidation of conventional Ni3Al-based alloy in isothermal conditions forms a double-layered scale, consisting of an outer NiO and inner Al2O3 as the oxide layer in Fig. 8. From the aforementioned results, thermal expansion and contraction of the alloy lead to cracks at the oxide layer/ matrix interface during cooling process. During the long-term cyclic oxidation process, local damage of the outer layer happened and on the 10

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Fig. 11. Schematic diagram of the overall description of transient oxidation.

conclusions are proved by surface scanning analysis (Fig. 12c). Al2O3 dominates the oxidation process due to its faster growth rate. In addition, insufficient oxygen combines with the Ni in the matrix, so that the presence of NiO is hardly observed in the oxide layer. It is worth noting that with continuous cyclic oxidation the outer oxide layer is substantially exfoliated due to the difference in thermal expansion coefficient during multiple cooling (Fig. 12a). When the alloy is oxidized at 1100 °C for 80 h, the O atoms in the air continuously diffuse into the matrix to increase the thickness of the Al2O3 oxide layer which is about 3 μm. The increase of the cyclic oxidation time causes that Ni diffuses to the surface of the specimen to form a small amount of NiO (Fig. 12b). According to Fig. 10 (a), some pores are generated on the surface after the oxide scales peeled off, and the formation of these pores is considered to be caused by the diffusion of alloying elements [58]. Pores would form at the oxide layer/substrate interface as inward diffusion of Ni is faster than outward diffusion of Al [58]. According to Fig. 10 (e) and 12, after a long period of oxidation, oxide scale is mainly composed of protective dense Al2O3 without more voids, which indicates that the diffusion of Ni is hindered, meaning thickening of Al2O3 layer also hinders the outward diffusion of the Ni, which is one of the important reasons for the thinning of outer oxide layer.

surface area of the damaged scale, O atoms are more likely to diffuse, causing an increase in oxygen partial pressure on the damaged area of the scale. At the same time, due to the formation of Al2O3, the Al element is consumed, and the aluminum content is lower than the content of Ni. Therefore, NiO first generated in the damaged area of the scale instead of Al2O3. Rapid growth dominates by outward diffusion of Ni cations allows oxides to grow outward through cracks, thus the spherical NiO forms. With the proceeding of cyclic oxidation, Ni cations continue to diffuse towards the alloy surface, resulting in the growth of spherical NiO. In addition, due to the relatively fast penetration of oxygen through the scale, oxides form rapidly at the inner oxide layer [44]. Due to the thermo dynamical stability of Al2O3, a little oxygen is sufficient to form this phase and the most stable phase present in the inner oxide layer is Al2O3 [44]. Therefore, it is the beginning of Al2O3 layer formation. At early stage of oxidation, after all Al elements from a certain volume increment is consumed, enough oxygen left form less stable oxides (NiAl2O4, Cr2O3 and NiCr2O3). The polycrystalline Ni3Albased superalloy after long-term cyclic oxidation, due to the continuous growth and spallation behavior of the scale, has a large amount of pores existing on the surface of the specimen and loosing structure. At this time, O in the air will continuously diffuse along the oxide pores to the matrix, resulting in thickening of the Al2O3 layer [22]. The distribution of the oxide layer after long-term cyclic oxidation is shown in Fig. 12. With cyclic oxidation time of 40 h, a layer of Al2O3 with dimensions of about 2 μm forms on the surface of specimen (Fig. 12a). The above

5. Conclusions Oxidation behaviors of a polycrystalline Ni3Al-based alloy under

Fig. 12. Secondary SEM images of cross-section morphology of solution treated alloy after cyclic oxidation for: (a) 40 h, (b) 80 h, (c) surface scanning results. 11

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different conditions (as-cast and solution treated) were investigated and the following conclusions were obtained. For the as-cast alloy, due to the lower Al content, higher internal energy and crystal defects around the (γ-γ′)/(γ+γ′) interfaces, oxidation initially occurred along the interface area and a large amount of NiO particles formed during the initial stage of cyclic oxidation at 1100 °C. Meanwhile, during the shortterm cyclic oxidation process, the increase of oxygen partial pressure and insufficiency of Al content lead to the reformation of NiO on the oxides peeling areas. While, for the solution treated alloy, at the initial stage of cyclic oxidation apart from NiO discontinuous Al2O3 appeared around the interface areas due to the increase of Al content. Meanwhile, during the cyclic oxidation process O will continuously diffuse into the matrix along the pores of the oxides which are caused by the continuous growth and spalling of the scales, resulting in the thickening of Al2O3 layer which is beneficial for improving the oxidation resistance of the solution treated alloy.

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