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Jiu-han Xiao, Ying Xiong, Li Wang, Xiang-wei Jiang, Dong Wang, Kai-wen Li, Jia-sheng Dong, and Lang-hong Lou, Oxidation behavior of high Hf nickel-based superalloy in air at 900, 1000 and 1100°C, Int. J. Miner. Metall. Mater., 28(2021), No. 12, pp.1957-1965. https://dx.doi.org/10.1007/s12613-020-2204-z
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To investigate the oxidation behavior of a nickel-based superalloy with high hafnium content (1.34wt%), this study performed isothermal oxidation tests at 900, 1000, and 1100°C for up to 200 h. X-ray diffraction and scanning electron microscopy with energy-dispersive X-ray spectroscopy were applied to study the oxidation behavior. The weight gain of the high Hf nickel-based superalloy exhibited a parabola-like curve, and no spallation of the oxide scale was observed during the oxidation tests. The alloy presented excellent oxidation resistance, and no HfO2 was observed in the oxide scale at 900°C. With the increase of the oxidation temperature to 1000°C, HfO2 particles formed in the spinel phases of the scale, and “peg-like” HfO2 was observed within and beneath the inner layer of Al2O3 after 200 h. As the oxidation temperature rose to 1100°C, “peg-like” HfO2 was observed at the early stage of the oxidation test (within 25 h). The formation mechanism of HfO2 and its impact on oxidation resistance were investigated based on the analysis of the oxidation test results at different temperatures.
Nickel-based superalloys have been widely applied to manufacture the blades and vanes of aero engines and gas turbines because of their outstanding oxidation resistance, superior corrosion resistance, and excellent mechanical capacities at high temperatures. For improving service temperatures, large amounts of refractory elements (such as W, Mo, Re, Hf, and Ta) have been mixed into newly developed superalloys. The large atomic radii of these refractory elements can enhance the solution strengthening effect for alloys. Moreover, the addition of refractory elements increases the melting points of alloys and reduces the cation diffusion ability. These effects can also improve mechanical properties at elevated temperatures. The influences of W, Mo, Ta, and Re on the performances of superalloys have been extensively investigated [1–4], but the effects of Hf have been rarely reported in the open literature [5–8].
In general, Hf is a grain boundary strengthening element that forms dispersed HfC along grain boundaries. HfC can prevent the movement of dislocation and, thus, pin the grain boundaries to keep them from sliding, which greatly improves the creep resistance of superalloys at high temperatures. The hardness and liquidity of superalloys can also be increased remarkably by HfC. Therefore, high Hf content helps improve the service and casting properties of superalloys [9–10]. However, Hf also exerts detrimental effects on Ni-based superalloys. For example, the interface reaction between molten alloy and ceramic shell at high temperatures can be promoted with the addition of Hf, thereby increasing the risk of the formation of surface defects and impurities [11–13]. At present, most researches on the effects of Hf is mainly focused on the oxidation resistance and adherence of coatings and, thus, few research explored the high Hf content of alloys [14–19]. Hf is widely known to benefit oxidation resistance. Remarkable improvements in the adherence of oxide scales have been made through the addition of Hf to coatings [14−16]. However, limited studies have investigated the effects of high Hf content on the oxidation resistance of Ni-based superalloys. Related studies have reported that the quantity of Hf in commercial nickel-based superalloys is relatively low (usually 0 or <0.4wt%) [20−21]. Additionally, the effects of high Hf content on oxidation resistance are not systematically investigated. To evaluate the oxidation behavior of high hafnium superalloys, the current work conducted oxidation tests on a Ni-based superalloy with 1.34wt% Hf at 900, 1000, and 1100°C. The oxidation behavior of a Ni-based superalloy with high Hf content was evaluated and discussed accordingly.
The alloy was prepared in a vacuum induction melting furnace and casted into cylindrical rods by pouring the liquid alloy into a ceramic shell mold prepared by the lost wax technique. The chemical composition of the high Hf nickel-based superalloy is listed in Table 1. In order to eliminate casting defects, reasonable melting parameters should be adopted. The shell mold was prebaked at 1100°C for 4 h. Meanwhile, the ingot was heated up to 1420°C and held for 10 min until complete melting was observed. The alloy was solution heat treated at 1185°C for 2 h, air cooled, and then aged including two steps: at 1050°C for 4 h, followed by air cooling, and then at 870°C for 20 h, followed by air cooling. The oxidation specimens measuring 10 mm × 20 mm × 1.5 mm were cut with an electro-discharge machine. All the specimens were grounded with 1000-grit sandpaper under water and then ultrasonically cleaned in acetone for 5 min before the oxidation tests. At the same time, corundum crucibles were presintered at 1200°C for 100 h.
| Ni | Co | Cr | Al + Ti | W + Mo | Ta | Hf | C | B | Zr |
| Bal. | 10.02 | 8.37 | 6.42 | 10.60 | 3.03 | 1.34 | 0.13 | 0.02 | 0.01 |
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Oxidation tests were performed at 900, 1000, and 1100°C in accordance with the service temperature of the high Hf nickel-based superalloy for up to 200 h in air. The weight gain of the specimens was measured after 25, 50, 75, 100, 125, 150, 175, and 200 h, and the typical specimens (oxidation interrupted after 1, 5, 25, 100, and 200 h) were selected for analysis. The test specimens were weighted with an analytical balance with a sensitivity of 10−4 g to ensure the accuracy of the experiment. The average value of the weight gained from three equivalent specimens in this study was calculated to reduce error propagation under the same conditions. The weighed data were unified into the unit of mg/cm2 to confirm the specific weight change.
The oxidation products were analyzed by X-ray diffraction (XRD). The surface and cross sections of the oxidized specimens were observed and analyzed by scanning electron microscopy (SEM) equipped with energy-dispersive X-ray spectroscopy and backscatter detectors.
The curves of weight gain as a function of thermal exposure time for the high Hf nickel-based superalloy at 900, 1000, and 1100°C are plotted in Fig. 1. Weight gain throughout the whole oxidation process at different temperatures was observed, and the weight increment increased with the oxidation temperature. The oxidation kinetics curve at each oxidation temperature roughly followed the parabolic law. As the oxidation temperature rose, the rate of initial oxidation increased, and the weight gain of oxidation became increasingly obvious.
The average weight gain of the high Hf nickel-based superalloy was about 1.95 mg/cm2 for 200 h at 1100°C. This value was about twice that obtained with isothermal oxidation at 1000°C (0.84 mg/cm2) and nine times that obtained with isothermal oxidation at 900°C (0.23 mg/cm2). The comparison of these results with the aforementioned sets of data reveals the obvious rule.
XRD analysis was performed to identify the oxidation products. Fig. 2 shows the XRD results of the original and oxidized specimens at 900, 1000, and 1100°C for 200 h. The diffraction spectrum shows that oxidation products scarcely formed at 900°C. At 1000°C, the oxidation products were mainly composed of Al2O3, HfO2, CrO, and two types of spinel phases ((Ni,Cr,Ti)(Cr,Co)2O4 and (Ta2O5)15(WO3)2). At 1100°C, the amount of oxidation products increased, and HfO2 was observed.
Fig. 3 shows the surface morphology of the high Hf nickel-based superalloy thermally exposed at 900, 1000, and 1100°C for 200 h. The specimen surfaces still presented some metallic gloss, and a few white lumpy particles were dispersed on the surface (Figs. 3(a)–3(c)). This result indicated that only a slight oxidation reaction occurred on the surface of the high Hf nickel-based superalloy at 900°C. At 1000°C, two types of surface morphology (light gray and dark gray) were observed (Figs. 3(d)–3(f)). This feature implied the occurrence of a more serious oxidation than that observed in the specimens oxidized at 900°C. Figs. 3(g)–3(i) show a highly complex and serious surface oxidation at 1100°C.
The cross sections of the oxidized specimens were observed. Fig. 4 shows the backscattered electron images and distribution maps of the major elements on the cross section of the high Hf nickel-based superalloy oxidized at 900°C for 200 h. An oxide layer formed on the surface of the specimen, and no γ′ depletion zone was found (Fig. 4(a)). Two typical regions, namely, noncarbide region and carbide region, were selected for measurement. A discontinuous scale rich in Al, Cr, and O was observed (Fig. 4(b)). These results, in combination with those of the XRD, indicate the formation of a discontinuous oxide layer composed of Al2O3 and CrO on the surface of the specimen in the noncarbide region. The thickness of the oxide scale was about 2.5 μm. However, in the carbide region, the oxidation process was accelerated by interdendritic MC carbides (Fig. 4(c)). Unlike noncarbide regions, the carbide region did not show an Al-rich surface. Al cations diffused along the interface of the carbides and matrix and reacted with the O anions to form Al2O3 surrounding the carbides. Meanwhile, the oxides of W, Ta, Hf, and Ti could be traced to the oxidized MC carbides. The depth of the inward diffusion of the O anions reached 14 μm. Owing to the excellent oxidation resistance of the alloy at 900°C, a slow weight gain was observed, and this condition led to a low-slope and flat oxidation kinetics curve (Fig. 1).
The cross sections of the oxide scales of the high Hf nickel-based superalloy oxidized at 1000°C are shown in Fig. 5. Only a thin oxide scale was observed after oxidation for 5 h. A light gray oxidation product wrapped in a grayish-white phase appeared in the oxide scale after 100 h of oxidation. Moreover, a dark gray layer continuously precipitated in the inner oxide scale. After oxidation for 200 h, the oxide scale deepened, and a fine strip-like bright phase surrounded by a dark gray phase precipitated in the matrix below the oxidation layer. A thin oxide layer (Al2O3) is usually observed around carbides because of the selective oxidation of Al [22–23]. Herein, a large number of bright white oxidation products were found in the vicinity of the oxidized MC carbide after oxidation for 200 h.
Fig. 6 shows the microstructure micrographs and element distribution of the cross section of the high hafnium superalloy oxidized at 1000°C for 200 h. The surface of the specimen was covered by a continuous oxide layer, and an obvious γ′ depletion zone appeared under the oxide scale (Fig. 6(a)). The thickness of the γ′ depletion zone was about 15 μm. Some of the interdendritic MC carbides exposed on the surface were oxidized; they are marked as yellow arrows in Fig. 6(a). As shown in Fig. 6(b), the oxide scale was divided into the outer and inner layers. The outer oxide layer was complex and mainly composed of a Ni, Co, Cr, and Ti-rich phase, W and Ta-rich phase, and Hf-rich oxide. Based on these results, along with the XRD results, the oxidation products were determined as (Ni,Cr,Ti)(Cr,Co)2O4, (Ta2O5)15(WO3)2, and HfO2. The dark gray phase located in the inner oxide layer was identified as Al2O3. The bright white phase embedded in Al2O3 was confirmed as HfO2. In the carbide region, the blocky MC carbides underwent an oxidation reaction and formed HfO2 in situ. Moreover, an Al2O3 layer was found on the outer surface (Fig. 6(c)). The diffusion of Hf cations was more active during oxidation at 1000°C than during oxidation at 900°C. Additionally, HfO2 was formed within the outer oxide layer and beneath the inner layer. The oxidation of the alloy was aggravated, and the protective oxide layer was established gradually at this moment, resulting in a smooth parabola of the oxidation kinetics curve (Fig. 1).
The microstructure of the high hafnium superalloy after oxidation at 1100°C is shown in Fig. 7. The oxide scale of about 5 μm was formed after 1 h of oxidization. The thickness of the oxide scale increased after oxidation for 5 h. A few bright white particles (HfO2) precipitated beneath the inner oxide layer after oxidation for 25 h. The amount of HfO2 that precipitated under the oxide scale increased with the increase of the oxidation time to 200 h. The precipitated HfO2 phase tended to extend inward in the direction perpendicular to the specimen surface. Serious oxidation occurred along the grain boundary, and the depth of the oxidized grain boundary reached 50 μm (Fig. 7(f)).
Fig. 8 shows the element distribution maps of the cross section of the high Hf nickel-based superalloy after oxidation at 1100°C for 200 h. A limited number of oxidized interdendritic carbides were observed on the surface, and a relatively thick γʹ depletion zone (approximately 40 μm) was found beneath the oxide layer (Fig. 8(a)). Therefore, the oxidation at 1100°C was dominated by uniform diffusion, and a large number of Al atoms in the γʹ phase diffused to the surface oxide scale. Compared with the oxidation at 1000°C, the oxidation at 1100°C presented a great increase in oxide thickness and resulted in the embedding of many strip-like HfO2 surrounded by Al2O3 in the matrix. This result implied that Hf cations were more active and diffused much faster at 1100°C than at 1000°C. Before the first 25 h of oxidation at 1100°C, the stable protective layer was not established completely, thereby leading to the significant weight gain of the alloy. Subsequently, the weight gain increased slowly when the stable oxide scale on the matrix formed. The corresponding oxidation kinetics curve is shown in Fig. 1.
The experimental results clearly showed that the addition of Hf element exerted different effects on the oxidation behavior of the high Hf nickel-based superalloy at different temperature. For isothermal oxidation at 900°C, the Hf cations showed no obvious diffusion, but the MC carbides showed oxidation. On the one hand, this result is attributed to the low thermodynamic driving force and high barrier energy for the diffusion of Hf cations at a relatively low temperature. According to the recent literature and thermodynamic software calculations (Outotec HSC Chemistry), the equilibrium Gibbs free energy of the reactions shown in Eqs. (1)–(3) are negative at 900°C [24–26].
| 43Al+O2=23Al2O3 | (1) |
| 4Cr+5O2=2Cr2O5 | (2) |
| Hf+O2=HfO2 | (3) |
Hence, the reactions to form Al2O3, CrO, and HfO2 can theoretically occur. However, because the cationic radius of Hf is larger than those of Cr and Al, the diffusion barrier energy of Hf cations is higher than those of Al and Cr cations. This characteristic leads to a faster outward diffusion of Al and Cr cations than Hf cations from an oxide layer. The oxide scale of Al and Cr on the surface can prohibit the further diffusion of oxygen anions into the metal matrix to generate HfO2 [27–28]. However, the atomic arrangement of MC carbides that precipitated in the interdendrite during the solidification process was different from that of the surrounding matrix. This lattice mismatch caused the energy to rise at the carbide/matrix interface [29]. Hence, oxygen could diffuse rapidly inward to oxidize MC carbides through the high energy interface of the carbide/matrix. At this point, quantities of active Al cations in the γʹ phase also diffused to the carbide/matrix interface and combined with oxygen to form Al2O3. As a result, the Al2O3 layer on the surface disappeared in the carbide region.
The mutual diffusion between the reactive element cations and the oxygen anions became more serious at 1000°C than at 900°C. The diffusion coefficient of Hf increases at a high oxidation temperature according to the following Arrhenius formula [30]:
| D=D0exp(−QRT) | (4) |
where D is the diffusion coefficient, D0 is the pre-exponential factor, Q is the diffusion activation energy, R is the gas constant, and T is the diffusion temperature. In this formula, D0 and Q are constants related to alloy composition. R is a constant equal to 8.314 J/(mol·K). Thus, diffusion coefficient D increases as temperature T increases. In addition to the significant outward diffusion and local segregation of Hf cations within and beneath the oxide scale, blocky MC carbides were completely oxidized for HfO2 formation. HfO2 was dispersedly distributed in the outer spinel phase within 100 h. Previous studies have suggested that reactive elements (such as Hf) are prone to segregation at grain boundaries and to reaction with oxygen anions that diffuse inward along the grain boundaries to form HfO2 [31–36]. The latter view was contested by Cao et al., who reported the oxidation mechanism of HfO2 particles in the spinel, specifically in this mechanism, Hf cations dissolved in the matrix are oxidized in situ by oxygen anions that diffuse inward, and the continuous outward diffusion of Hf cations in the matrix driven by the oxygen potential gradient make HfO2 particles coarse [37–39]. With respect to the bright white HfO2 phase that resembled “pegs” beneath the inner layer of Al2O3 in this work, the formation mechanism is different from that in spinel. As a result of the existence of an external diffusion driving force, the Hf cations dissolved in the matrix easily diffused outward for outward diffusion through high energy defects, such as vacancies, interfaces, and grain boundaries. This action led to the segregation of Hf cations at the high energy defect sites. Meanwhile, the O anions diffused inward rapidly along the grain boundaries in the oxide scale and combined with the Hf cations that segregated in the defects to form HfO2. As for HfO2, the conductivity of oxygen anions was higher than that of Al2O3 because of their defective crystal structure [39–40]. The O anions in HfO2 easily reacted with the Al cations in the surrounding matrix to form an Al2O3 layer. The HfO2 that formed in the defects acted as a core that continued to grow inward and subsequently exhibited a “peg-like” shape. Moreover, the carbides on the surface and subsurface were preferentially oxidized at 1000°C. The oxidation mechanism of the carbides on the surface at 1000°C was the same as that at 900°C. However, for the oxidized MC carbides on the subsurface, the principle of oxidation is attributable to the rapid inward diffusion of O anions to the carbide/matrix interface through the short-circuit diffusion of grain boundaries in the oxide scale, resulting in the reaction with Hf cations in carbides to form HfO2. The HfO2 that formed on the outer surface of the carbides acted as a fast transport medium for O anions, thereby continuously transforming the internal HfC into HfO2.
Apart from the oxidation behavior at 900 and 1000°C, relatively small number of oxidized MC carbides and large amounts of “peg-like” HfO2 appeared near the surface after oxidation at 1100°C. This result indicated the increasingly active diffusion of the Hf cations dissolved in the matrix. The effect is attributed to the increase in the diffusion driving force of the reactive elements caused by a rise in temperature. Hf cations can easily overcome energy barriers to achieve diffusion. The thickening of the γʹ depletion zone and the formation of Kirkendall pores are strong evidence of vigorous mutual diffusion [41]. More Hf cations dissolved at a deeper location of the matrix diffused outward promptly and segregated at high energy sites to form a “peg-like” HfO2 cluster. The reduction of the oxidized surface of the MC carbides is probably due to the decomposition of some MC carbides into M6C during the long-term isothermal exposure [42]. As O anions could not readily diffuse out of the dense Al2O3 layer wrapped by “peg-like” HfO2, the carbides on the subsurface were hardly oxidized. A serious erosion of grain boundaries was also observed. As a result of the extremely high level of grain boundary energy, the grain boundary became a fast channel for oxygen anion diffusion. Vulnerable grain boundaries that are attacked during service may trigger crack initiation.
Based on the investigation of the effects of Hf on the oxidation behavior of the high Hf nickel-based superalloy at different temperatures, the following conclusions can be drawn. (1) In contrast to the oxidation at 900°C, HfO2 particles formed within and beneath the oxide layer at 1000 and 1100°C. With increasing temperature, large amounts of “peg-like” HfO2 formed. (2) HfO2 formed as follows: Hf dissolved in the matrix was precipitated in situ, HfO2 transformed from HfC, and “peg-like” HfO2 was formed by diffusion. (3) The grain boundary can be used as a short-circuit channel of diffusion for oxygen anions and reactive element cations. This application increases the susceptibility to erosion at high temperatures.
This work was financially supported by the National Science and Technology Major Project (Nos. 2017-VI-0001-0070 and 2017-VI-0019-0091) and the National Natural Science Foundation of China (Nos. 51631008, 51771204, U1732131, 51671196, and 91860201).
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| Ni | Co | Cr | Al + Ti | W + Mo | Ta | Hf | C | B | Zr |
| Bal. | 10.02 | 8.37 | 6.42 | 10.60 | 3.03 | 1.34 | 0.13 | 0.02 | 0.01 |
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