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Article

Effect of Annealing Temperatures on Phase Stability, Mechanical Properties, and High-Temperature Steam Corrosion Resistance of (FeNi)67Cr15Mn10Al5Ti3 Alloy

by
Zhen Cai
1,
Fengyang Jiang
2,
Na Wei
3,4,*,
Lei Mi
3,4,
Chenhui Zhang
3,4,
Xiaohua Liu
3,4,
Fang Si
2 and
Tiandong Wu
3,4
1
Power Station Life Management Technology Center, Suzhou Nuclear Power Research Institute, Suzhou 215004, China
2
School of Materials Engineering, Xi’an Polytechnic University, Xi’an 710048, China
3
Technology Center, Xi’an Chao Jing Technology Co., Ltd., Xi’an 710200, China
4
Shanxi Key Laboratory of Advanced Metal Structural Materials Precision Thermoforming, Xi’an 710200, China
*
Author to whom correspondence should be addressed.
Metals 2022, 12(9), 1467; https://doi.org/10.3390/met12091467
Submission received: 1 July 2022 / Revised: 23 August 2022 / Accepted: 23 August 2022 / Published: 1 September 2022
(This article belongs to the Special Issue High-Entropy Alloys: Structures, Properties and Applications)
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Versions Notes

Abstract

:
The effect of different annealing temperatures on the phase stability and mechanical properties of (FeNi)67Cr15Mn10Al5Ti3 high-entropy alloys (HEAs) was studied. The phase stability was analyzed by X-ray diffraction (XRD), scanning electron microscopy (SEM), and electron backscattering diffraction (EBSD). The mechanical properties of the alloy were characterized by hardness and tensile tests. Furthermore, the heat-resistant corrosion properties of the (FeNi)67Cr15Mn10Al5Ti3 alloy after annealing at 800 °C was tested under high-temperature steam. The results indicated that HEAs exposed to different annealing temperatures always exhibited the face-centered cubic (FCC) phase. With rising annealing temperature, the dendrite structure of the alloys in the as-cast condition gradually disappeared, with recrystallization and precipitation of larger grains. The tensile strength of the alloy first increased and then decreased with the rising annealing temperature, the hardness and yield strength of the alloy decreased slightly, and the tensile elongation varied greatly. These findings can be used as a basis for improving the phase stability and mechanical properties of a Cr-Fe-Ni-Mn-HEA system with unequal atomic ratios. The heat and corrosion resistance of the alloy at 360 °C and 400 °C was better than that of Zr-4 alloy.

1. Introduction

Much attention has been paid to high-entropy alloys (HEAs) because of their excellent properties and wide range of possible applications [1,2,3,4,5,6]. In HEA systems, HEAs with equal atomic ratios have been frequently studied because of their typical structure, where a single solid solution phase is formed [7,8,9,10,11]. Annealing is considered as a simple and effective method to improve the microstructure and mechanical properties of HEA. However, microstructure and composition analyses at different scales indicate that the HEA’s solid solution is unstable when annealed at lower annealing temperatures. For example, when exposed to temperatures between 600 °C and 800 °C, isoatomic CrMnFeCoNi alloys exhibited tetragonal σ phases, and all the elements were observed at the grain boundary of the FCC phase. Such unstable phases may result in changeable mechanical properties [12,13,14,15]. Therefore, the phase stability of HEA requires further consideration [16,17].
Alloy components are essential to the phase formation and stability of HEA [18,19,20]. For example, the σ phase can be formed easily by changing the component concentration. Some studies [21,22,23,24] on Cr-Mn-Fe-Co-Ni alloys with unequal atomic ratios show that these alloys have a lower fault energy and more easily form the σ phase, which leads to decreased toughness. To ensure a stable phase structure, no new precipitated phase can be found in the solution of trace elements, and the structure type of the material should be unchanged. For example, Al is usually used as the BCC stabilizer of the AlxCoCrFeNi alloy, whereas Co and Ni are the FCC stabilizers of the CoCrFeNiCuAl alloy [18,19,25,26,27]. Many studies have focused on the structure and properties of HEA with equal atomic ratios. However, the effect of annealing treatment on the microstructure and mechanical properties of HEA with unequal atomic ratios under the action of trace elements remains unclear [4,22,23,24]. In order to cope with complex environments such as high temperature, corrosion, and vibration, the alloy should not only have good mechanical properties, but also have excellent heat and corrosion resistance. The (FeNi)67Cr15Mn10Al5Ti3 high entropy alloy contains high levels of the anti-oxidation and anti-corrosion elements Cr and Al, so its thermal corrosion resistance was studied in order to explore the possibility of replacing the general Zr-4 alloy, which is a good cladding material in the nuclear field.
In this study, microalloyed (FeNi)67Cr15Mn10Al5Ti3 alloys with unequal atomic ratios were explored to investigate the effects of different annealing temperatures (700–1000 °C) on the phase stability and mechanical properties of the alloy and its heat and corrosion resistance compared with Zr-4 alloy. The results contribute to the understanding of the phase stability of HEA systems with unequal atomic ratios and to adjust their microstructure and mechanical properties, laying a theoretical basis for the further design and application of high-strength HEA.

2. Materials and Methods

The following metallic materials with different mass ratios of Fe (34.4 wt.%), Cr (14.3 wt.%), Mn (10.1 wt.%), Ni (36.1 wt.%), Al (2.47 wt.%), and Ti (2.63 wt.%) were placed in a vacuum induction furnace for melting (Xi’an Chao Jing Technology Co. Ltd, Xi’an, China). After this, 5 wt.% Mn was utilized for vacuum arc re-melting four or five times. After cooling for 30 min, the (FeNi)67Cr15Mn10Al5Ti3 HEA in the as-cast condition was obtained using a water-cooled metal mold. The as-cast HEA was annealed at 700 °C, 800 °C, 900 °C, and 1000 °C with the same heating time in a resistance furnace for 2 h before air cooling.
After mechanical polishing and etching, the annealed samples were analyzed by a Shimadzu 6000X X-ray diffractometer (Shimadzu Co., Ltd., Zhongjing District, Kyoto, Japan) and a JSM-7900F field emission scanning electron microscope (Japan Electronics Co., Ltd., Zhaodao, Tokyo, Japan) for grain size and phase structure. The polished samples were etched for 20–50 s in aqua regia. The microstructure and composition of the etched samples were then observed with a FEI Quanta 400F scanning electron microscope (FEI Co., Hillsborough, OR, USA). At room temperature, the tensile properties of the samples were tested using a JX-50A electronic universal material testing machine (Shen Zhen Xin Aansi material testing Co., Ltd., Shenzhen, China).
The (FeNi)67Cr15Mn10Al5Ti3 alloy after heat treatment at 800 °C was subjected to a corrosion test in an autoclave. The corrosion solution was 1000 ppm boric acid and 2 ppm lithium hydroxide, and the temperature and pressure were 360 °C, 18.7 Mpa, 400 °C, and 10.3 MPa respectively. The sampling times were 3 days, 10 days, and 14 days. Each sampling was weighed using an electronic balance, and the corrosion morphology was observed and analyzed with a scanning electron microscope.

3. Results and Discussion

3.1. Analysis of Phase and Microstructure Evolution

Figure 1 shows the XRD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA at different treatment temperatures. Figure 1a shows the FCC phase of the alloy in the as-cast condition. Studies on the formation of a solid solution of HEA showed that a simple solid solution was formed only when the competition parameter of the mixing enthalpy and mixing entropy of HEA was ω ≥ 1.1 and the atomic radius difference was δ ≤ 6.6% [28,29,30,31,32]. Moreover, the mixing entropy and mixing enthalpy of (FeNi)67Cr15Mn10Al5Ti3 alloy were 12.49J/mol/K and −9.61kJ/mol, respectively, the competitive parameter ω was 1.9, and the atomic radius difference δ was 4.67%. Therefore, the parameter values of (FeNi)67Cr15Mn10Al5Ti3 HEA fell into this range, which was consistent with the test results. No other phase was observed with increasing annealing temperature, which indicated that the HEA had a good phase stability. High mixing entropy promotes the mutual dissolution of elements and inhibits the appearance of ordered phases such as intermetallic compounds [33,34]. Meanwhile, the stable phase is also guaranteed by the design of unequal principal elements in the high Ni alloy.
The peak offset was found, as shown in Figure 1b, in the (111) enlarged view of the crystal plane, which indicated that the lattice constant calculated by indexation (a = b = c = 0.3063 nm) was different from the standard constant (a = b = c = 0.3592 nm). The change in the X-ray diffraction angle for (111) suggests a change in the lattice constant. Obviously, lattice distortion occurred, which might have been caused by different elements, aggregated by segregation, occupying the FCC point locations. According to the XRD patterns, the diffraction peak area became larger after the annealing process, which can be explained by the existence of diffusion and the growing solid solution during annealing. The higher temperature led to an enhanced diffusion effect in the alloying elements, thus forming more fine grains. Moreover, the melting of the primary phase led to grain refinement, which then increased with the prolongation of the holding time.
Figure 2 shows the microstructure morphology and point energy spectrum analysis of (FeNi)67Cr15Mn10Al5Ti3 HEA in its as-cast condition. Figure 2a,b show that the as-cast (FeNi)67Cr15Mn10Al5Ti3 exhibited a typical dendrite structure, consisting of uneven black dendrite structures (DR) and interdendrite structures (ID). According to Figure 2b, the ID structure was loose, and there were interdendrite gaps. The point energy spectrum test results at A and B in Figure 2b are shown in Table 1. The Al and Ti of the as-cast sample segregated at the ID for different radii, valence electron concentrations, and electronegativities. This means that the DR of the as-cast HEA was an FCC structure enriched with the Cr element, and the ID of the as-cast HEA was an FCC structure enriched with Al and Ti elements, which resulted in the different brightnesses between the DR and ID regions.
The microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures is shown in Figure 3a–h. The dendrite structure of the as-cast alloy gradually became uniform, and the ID structure decreased with the increase in the annealing temperature, and the microstructure was divided into two parts with different bright and shaded areas, as shown in Figure 3a,c,e,g. The ID was a loose microstructure and the DR was dense with a rhomboid black isolated precipitated phase (marked by the circle) from the high-magnification topography shown in Figure 3b,d,f,h. Combined with the distribution of elements in the as-cast alloy structure, the dark part, including the isolated precipitated phase, was the solid solution of the DR with a Cr-rich FCC structure, and the bright area was the solid solution of the ID with a Al- and Ti-rich FCC structure. The formation of such a microstructure was mainly caused by the high entropy effect. When the annealing temperature was lower than 800 °C, the separated elements diffused and tended to become uniform, resulting in the gradual reduction in the Al- and Ti-rich FCC structure phase. However, when the temperature was higher than 800 °C, the Al- and Ti-rich FCC structure phase formed easily due to the mixing entropy of the high-entropy alloy inhibiting the diffusion of the segregation elements.
Figure 4 shows the EBSD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures. Figure 4a to 4e illustrate that the grain size varied greatly, and the grain size became heterogeneous and then gradually became uniform with the increasing annealing temperature, as shown in Figure 4f. The average grain size was 13.2 μm, and a clear grain boundary was seen when the HEA was annealed at 800 °C. At 1000 °C, the crystal grain size was relatively uniform, the crystal grains became round, and the average crystal grain size was about 23.8 μm. Therefore, at a low annealing temperature (≤800 °C), the elements of (FeNi)67Cr15Mn10Al5Ti3 HEA diffused through the grain boundary, and the solid solution became larger and precipitated out. The microstructure gradually became uniform, but with a heterogeneous grain size distribution as the precipitation time changed. At a higher annealing temperature (>800 °C), the mixing entropy of the alloying elements increased, diffusion at grain boundaries was suppressed, and the unevenly distributed elements presented with a uniform grain size.

3.2. Hardness and Tensile Properties

Figure 5 shows the hardness change in (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures. The hardness of the alloy declined from 327.72 HV for the as-cast condition to 240.33 HV at 900 °C, decreasing by about 26.7%. When exposed to an annealing temperature of less than 800 °C, the solid solution structure homogenized, and grain coarsening led to a decrease in hardness [35]. At 1000 °C, the hardness decreased by 10.4% compared with that of the as-cast alloy. The solid solution simultaneously precipitated and increased at high temperature, with a small grain size (as shown in Figure 3h), and the hardening effect of the fine grain increased the hardness of the alloy.
Figure 6 and Table 2 show the stress–strain curves and tensile properties of (FeNi)67Cr15Mn10Al5Ti3 HEA in the as-cast and annealed conditions at different annealing temperatures. The tensile strength and yield strength first increased and then decreased, rising from 821.3 MPa to 919.2 MPa. The fracture structure was obviously a cleavage fracture, with a few dimples, and the alloy was a brittle fracture. When the annealing temperature was ≥800 °C, the tensile strength of the alloy fell into the range of 645–689 MPa, decreasing by 16–21%. The fracture structure was mainly dimpled with good ductility, which was consistent with the change in hardness [36]. In contrast, the tensile elongation increased significantly from about 29.1% to 50–72%. The decline in yield strength and the increase in tensile elongation were caused by homogenization of the single solid solution and coarsening of the grain size. At a low temperature, the homogenization of the composition and the grains, without any growth, led to the enhancement of strength. However, the hardness decreased due to the elimination of the residual stress from casting. Although the homogenization of the composition became complete as the temperature increased, the grains grew, which led to the decrease in a strength and hardness. This also indicates that changing the annealing temperature does not precipitate the brittle σ phase, but changes the solid solution uniformity and grain size. These findings are consistent with the results of Zhu [37] and Salischev [38].

3.3. Hot Corrosion Properties

3.3.1. Corrosion Weight Loss Curve

Figure 7 shows the weight loss curves of corrosion samples of the (FeNi)67Cr15Mn10Al5Ti3 alloy and Zr-4 alloy at 360 °C and 400 °C, respectively. It can be seen that both the (FeNi)67Cr15Mn10Al5Ti3 and Zr-4 alloys gained weight in a 360 °C high-temperature and high-pressure water environment. With the increase in corrosion time, the weight gain of (FeNi)67Cr15Mn10Al5Ti3 increased, but it was lower than that of the Zr-4 alloy. In a 400 °C high-temperature steam environment, both the (FeNi)67Cr15Mn10Al5Ti3 and Zr-4 alloy gained weight, and with the increase in time, the weight gain degree of (FeNi)67Cr15Mn10Al5Ti3 was significantly lower than that of the conventional Zr-4 alloy.

3.3.2. Hot Corrosion Morphology

Figure 8 shows the microstructure of the (FeNi)67Cr15Mn10Al5Ti3 and Zr-4 alloys after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days. After 14 days of corrosion in a 360 °C high-temperature and high-pressure water environment, the oxide particles on the surface of (FeNi)67Cr15Mn10Al5Ti3 were unevenly distributed, and there were coarse particle areas and fine particle areas. By scanning the composition of O, Fe, Cr, Mn, Ni, Al, and Ti at low magnification, it was seen that the poor Cr, Ni, and Ti were obviously enriched in the coarse area of the oxide particles (Figure 9). After 14 days of corrosion in a 360 °C high-temperature and high-pressure water environment, there were no obvious corrosion products on the surface of the Zr alloy. By scanning the surface of the O, Zr, and Sn components, it was seen that the components were evenly distributed and there was no element aggregation (Figure 10). The composition analysis showed that the content of O was 29.4 wt.%, Zr was 69.3 wt.%, SN was 1.0 wt.%, and a small amount of Fe and Cr were observed at 0.2 wt% and 0.1 wt.%, respectively.
Table 3 shows the results of a point component analysis of Figure 11. It can be seen from the point composition analysis that the overall Cr content in the large particle area was about 12%, which was about 15% lower than that observed in the small particle area. In the large particle area, the large particles were mainly oxide with a low content of Ti. The bottom layer had a low oxygen content but with high contents of Ni, Mn, and Fe. The small particle area was covered by oxide, and the oxygen content was higher than that of the bottom layer. The small particle area had a high Cr content and low Mn, Ni, and Ti contents.
Figure 12 shows the microstructure of the (FeNi)67Cr15Mn10Al5Ti3 and Zr-4 alloys after corrosion in a 400 °C high-temperature and high-pressure water environment for 14 days. In the 400 °C high-temperature steam environment, the oxide particles on the surface of (FeNi)67Cr15Mn10Al5Ti3 were unevenly distributed, and there was a region without coarse oxide particles. The surface scan (Figure 13) of the O, Fe, Cr, Mn, Ni, Al and Ti components demonstrated that the oxygen content of the large particles was high and the oxygen content of base was low. In the area containing coarse oxide particles, the contents of Fe and Cr were high, the contents of Ni and Ti were low, and there was no significant difference between Mn and Al. In the 400 °C high-temperature steam environment, it can be seen from the Figure 14 that a large number of cluster products were distributed on the surface of the Zr alloy. The oxygen and chromium contents of these clusters were high and the Zr content was very low, which resulted from the oxidation of chromium. The point composition analysis of the clusters shows that the content of O was about 48.5 wt.%, the content of Cr was 42.2 wt.%, the content of Zr was 2.0 wt.%, and the rest were C elements. For the cluster free region, the content of O was about 27.4 wt.%, the content of Zr was 59.6 wt.%, the content of Sn was 1.0 wt.%, and the rest were C elements.

4. Conclusions

In this paper, the microstructure and phase stability of (FeNi)67Cr15Mn10Al5Ti3 HEA in as-cast conditions and at different annealing temperatures, and the change rules of alloy hardness and tensile properties were studied. The main conclusions are as follows:
The single FCC phase was found in the HEA in the as-cast and annealed conditions. The stable phase structure was consistent with the criterion of solid solution formation in HEA.
The transition of the solid solution microstructure of (FeNi)67Cr15Mn10Al5Ti3 HEA at the same annealing time was mainly caused by the solid solution and precipitation of Al and Ti in the FCC structure with the increasing temperature. The homogenization of solid solution was the main process, and the growth of grains was difficult, which resulted in a non-monotonous relationship between the strength and the hardness. However, at higher annealing temperatures, the diffusion of elements was inhibited, and the elements in the rich and poor areas of the solid solution were homogenized.
For unequal principal element (FeNi)67Cr15Mn10Al5Ti3 HEA, with an increase in the annealing temperature, a homogenized solid solution and a change in grain size were observed, which led to a change in the mechanical properties of the alloy. Low-temperature annealing can improve the tensile strength of HEA. At 700 °C, the tensile strength was 919.2 MPa, with an elongation of 26.3%. At higher annealing temperatures, the ductility of HEA can be improved. When the annealing temperature was ≥800 °C, the elongation of the alloy more than doubled.
Both (FeNi)67Cr15Mn10Al5Ti3 HEA and Zr alloys gain weight in a 360 °C high-temperature and high-pressure water environment. With the increase in corrosion time, the weight gain of HEA increased, but it was lower than that of the Zr alloy. In a 400 °C high-temperature steam environment, both the (FeNi)67Cr15Mn10Al5Ti3 and zirconium alloys gained weight, and with the increase in time, the weight gain degree of (FeNi)67Cr15Mn10Al5Ti3 was significantly lower than that of the conventional Zr alloy.

Author Contributions

Conceptualization, Z.C.; methodology, F.J.; validation, N.W. and L.M.; formal analysis, N.W.; investigation, C.Z.; resources, X.L.; data curation, F.S.; writing—original draft preparation, F.J; writing—review and editing, F.S.; project administration, T.W. All authors have read and agreed to the published version of the manuscript.

Funding

Nationl Natural Science Foundation of China-Young Fund Project (51901019) “Cavitation water jet impact surface strengthening and stress corrosion mitigation of nuclear power welded joints”.

Institutional Review Board Statement

The study did not involve humans or animals.

Informed Consent Statement

Not applicable.

Data Availability Statement

All authors agree that all experimental data of this study can be published.

Acknowledgments

The authors are grateful for the support provided by National Natural Science Foundation of China-Young Fund Project (51901019) and Shanxi Key Laboratory of Advanced Metal Structural Materials Precision Thermoforming.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. XRD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA. (a) Diffraction pattern of 2θ between 20° and 100°. (b) (111) Enlarged view of crystal plane.
Figure 1. XRD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA. (a) Diffraction pattern of 2θ between 20° and 100°. (b) (111) Enlarged view of crystal plane.
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Figure 2. Microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA: (a) 500×; (b) 5000×; A and B is the point energy spectrum of pentagram marks.
Figure 2. Microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA: (a) 500×; (b) 5000×; A and B is the point energy spectrum of pentagram marks.
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Figure 3. Microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures: (a,b) 700 °C; (c,d) 800 °C; (e,f) 90 0°C; (g,h) 1000 °C.
Figure 3. Microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures: (a,b) 700 °C; (c,d) 800 °C; (e,f) 90 0°C; (g,h) 1000 °C.
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Figure 4. EBSD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA: (a) as-cast; (b) 700 °C; (c) 800 °C; (d) 900 °C; (e) 1000 °C; (f) grain size distributions.
Figure 4. EBSD patterns of (FeNi)67Cr15Mn10Al5Ti3 HEA: (a) as-cast; (b) 700 °C; (c) 800 °C; (d) 900 °C; (e) 1000 °C; (f) grain size distributions.
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Figure 5. Microhardness of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures.
Figure 5. Microhardness of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures.
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Figure 6. Stress–strain curves and microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures: (a) stress–strain curves; (b) as-cast; (c) 700 °C; (d) 800 °C; (e) 900 °C; (f) 1000 °C.
Figure 6. Stress–strain curves and microstructure morphology of (FeNi)67Cr15Mn10Al5Ti3 HEA at different annealing temperatures: (a) stress–strain curves; (b) as-cast; (c) 700 °C; (d) 800 °C; (e) 900 °C; (f) 1000 °C.
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Figure 7. Corrosion weight loss curve of the (FeNi)67Cr15Mn10Al5Ti3 alloy and Zr-4 alloy: (a) 360 °C; (b) 400 °C.
Figure 7. Corrosion weight loss curve of the (FeNi)67Cr15Mn10Al5Ti3 alloy and Zr-4 alloy: (a) 360 °C; (b) 400 °C.
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Figure 8. Microstructure after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) (FeNi)67Cr15Mn10Al5Ti3 alloy; (b) Zr-4 alloy.
Figure 8. Microstructure after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) (FeNi)67Cr15Mn10Al5Ti3 alloy; (b) Zr-4 alloy.
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Figure 9. Surface scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Fe; (c) Distribution diagram of Cr; (d) Distribution diagram of Mn; (e) Distribution diagram of Ni; (f) Distribution diagram of Al; (g) Distribution diagram of Ti.
Figure 9. Surface scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Fe; (c) Distribution diagram of Cr; (d) Distribution diagram of Mn; (e) Distribution diagram of Ni; (f) Distribution diagram of Al; (g) Distribution diagram of Ti.
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Figure 10. Surface scan of the Zr alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Zr; (c) Distribution diagram of Sn.
Figure 10. Surface scan of the Zr alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Zr; (c) Distribution diagram of Sn.
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Figure 11. Spot scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days.
Figure 11. Spot scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days.
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Figure 12. Microstructure after 14 days of corrosion in a 400 °C high-temperature and high-pressure water environment: (a) (FeNi)67Cr15Mn10Al5Ti3 alloy; (b) Zr-4 alloy.
Figure 12. Microstructure after 14 days of corrosion in a 400 °C high-temperature and high-pressure water environment: (a) (FeNi)67Cr15Mn10Al5Ti3 alloy; (b) Zr-4 alloy.
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Figure 13. Surface scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 400 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Fe; (c) Distribution diagram of Cr; (d) Distribution diagram of Mn; (e) Distribution diagram of Ni; (f) Distribution diagram of Al; (g) Distribution diagram of Ti.
Figure 13. Surface scan of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 400 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Fe; (c) Distribution diagram of Cr; (d) Distribution diagram of Mn; (e) Distribution diagram of Ni; (f) Distribution diagram of Al; (g) Distribution diagram of Ti.
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Figure 14. Surface scan of the Zr alloy after corrosion in a 400 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Zr; (c) Distribution diagram of Fe; (d) Distribution diagram of Cr.
Figure 14. Surface scan of the Zr alloy after corrosion in a 400 °C high-temperature and high-pressure water environment for 14 days: (a) Distribution diagram of O; (b) Distribution diagram of Zr; (c) Distribution diagram of Fe; (d) Distribution diagram of Cr.
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Table
Table 1. EDS results of microregion in (FeNi)67Cr15Mn10Al5Ti3 HEA.
Table 1. EDS results of microregion in (FeNi)67Cr15Mn10Al5Ti3 HEA.
ElementsAB
Al/wt.%3.512.62
Ti/wt.%6.743.50
Cr/wt.%7.9312.70
Mn/wt.%14.8411.06
Fe/wt.%18.7829.75
Ni/wt.%48.2040.36
Total100
Table
Table 2. Tensile properties of (FeNi)67Cr15Mn10Al5Ti3 HEA.
Table 2. Tensile properties of (FeNi)67Cr15Mn10Al5Ti3 HEA.
SamplesElongation after Breaking (%) Tensile Strength (MPa)Yield Strength (MPa)Specified Plastic Elongation Strength (MPa)
as-cast29.1821.3694.7570.9
700 °C26.3919.2777.9302.7
800 °C55.6645.2381.8190.9
900 °C50.3666.6402.2216.1
1000 °C72.6688.7363.0224.2
Table
Table 3. Spot scan composition of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days.
Table 3. Spot scan composition of the (FeNi)67Cr15Mn10Al5Ti3 alloy after corrosion in a 360 °C high-temperature and high-pressure water environment for 14 days.
Particles SizeAnalysis PointOFeCrMnNiAlTi
Large particle3640.017.512.86.520.21.71.3
Large particle4125.721.913.78.426.81.51.9
base3717.423.211.910.432.91.52.6
base4015.623.512.110.733.91.62.6
base4212.224.112.211.535.61.52.8
base4313.924.912.610.933.81.52.5
base4424.422.512.19.028.51.32.2
granule3824.124.415.47.925.11.61.5
granule3925.724.215.07.824.31.61.3
granule4529.523.014.57.322.81.51.4
granule4626.523.715.27.623.81.71.3
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Cai, Z.; Jiang, F.; Wei, N.; Mi, L.; Zhang, C.; Liu, X.; Si, F.; Wu, T. Effect of Annealing Temperatures on Phase Stability, Mechanical Properties, and High-Temperature Steam Corrosion Resistance of (FeNi)67Cr15Mn10Al5Ti3 Alloy. Metals 2022, 12, 1467. https://doi.org/10.3390/met12091467

AMA Style

Cai Z, Jiang F, Wei N, Mi L, Zhang C, Liu X, Si F, Wu T. Effect of Annealing Temperatures on Phase Stability, Mechanical Properties, and High-Temperature Steam Corrosion Resistance of (FeNi)67Cr15Mn10Al5Ti3 Alloy. Metals. 2022; 12(9):1467. https://doi.org/10.3390/met12091467

Chicago/Turabian Style

Cai, Zhen, Fengyang Jiang, Na Wei, Lei Mi, Chenhui Zhang, Xiaohua Liu, Fang Si, and Tiandong Wu. 2022. "Effect of Annealing Temperatures on Phase Stability, Mechanical Properties, and High-Temperature Steam Corrosion Resistance of (FeNi)67Cr15Mn10Al5Ti3 Alloy" Metals 12, no. 9: 1467. https://doi.org/10.3390/met12091467

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