The Effect of Heat Treatment on the Microstructure and Mechanical Properties of Al_0.6CoFeNi₂V_0.5 High Entropy Alloy

Abstract

Al_0.6CoFeNi₂V_0.5 high entropy alloy was successfully designed and prepared via the nonconsumable arc-melting process, and it was annealed at 600 °C, 800 °C, and 1000 °C for 4 h. Its microstructure and mechanical properties were studied. The as-cast alloy consisted of FCC and BCC phases, and no phase transformation occurred during annealing at 600 °C. Hard Al₃V-type metal compounds precipitated during annealing at 800 °C, and BCC particles precipitated in the FCC matrix during annealing at 1000 °C. After annealing, the strength and hardness of Al_0.6CoFeNi₂V_0.5 high-entropy alloy both showed a decreasing trend, because the annealing process eliminated the internal stress in this alloy. However, as the annealing temperature increased, the strength and hardness of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples gradually increased. This is because the hard Al₃V metal compounds precipitated when the annealing temperature was 800 °C, which produced the “second phase strengthening” effect. At 1000 °C, the larger volume fraction of the hard and fine BCC phase (21.81%) diffusely precipitated; the precipitation of this BCC phase not only produced a “second phase strengthening” effect, which also resulted in “solid solution strengthening”, ultimately exhibiting enhanced hardness and strength. These findings have important theoretical reference value for the study of the microstructure and mechanical properties of high-entropy alloys. And, this study plays a significant role in promoting the research and development of new component materials that bear compressive loads, such as columns in large factory buildings, supports for cranes, and clamping bolts for rolling mills in practical mechanical engineering.

1. Introduction

In actual mechanical engineering, many components, such as the columns of large factory buildings, the brackets of cranes, and the clamping bolts of steel rolling mills, bear compressive load during operation. Therefore, the development of new materials with excellent compressive mechanical properties is particularly important.

As a new type of alloy material, high-entropy alloys (HEAs) break through the design concept of traditional alloys, which comprises an alloy system with a composition number between 5 and 13, and the content of each main element is 5 at.%–35 at.% [1,2]. This new type of alloy extends the alloy design from the edge and corner regions of the phase diagram to the middle region. According to the high-entropy alloy design concept, if the toxic, radiative, and gaseous elements are removed and the remaining 75 stable elements are used as alloy components, the proportion of the alloy components composed of 3–6 elements reaches 219 million [3,4], greatly expanding the possibilities for alloy design. In traditional alloys, the tendency to form intermetallic compounds increases with the increase in the number of alloy element types. However, high-entropy alloys with equal or near equal atomic ratios tend to form simple solid solution structures in the as-cast state due to their high mixing entropy. And, they produce unique high-entropy effects, hysteresis diffusion effects, lattice distortion effects, and cocktail effects [5,6,7]; they exhibit high strength [8,9], high hardness [10,11], good wear and corrosion resistance [12,13], excellent high-temperature stability [14,15], and superconductivity at the macroscopic level [16,17]. As such, high-entropy alloys have great application prospects under extreme conditions and have become one of the research hotspots in the field of metal materials.

The study of the phase stability of high-entropy alloys under a wide temperature range and long-term service conditions would further promote their engineering applications. As such, researchers have conducted extensive studies on the influence of heat-treatment processes on the microstructure and properties of high-entropy alloys [17,18]. Jiang et al. [18] performed annealing at 1000 °C for 6 h of an AlCrFeNi₂Ti_0.5 HEA to promote the precipitation of some uniformly distributed fine particles inside the dendrites; moreover, the mechanical properties of the annealed AlCrFeNi₂Ti_0.5 alloy improved: a two-fold increase in plastic strain and an increase in compressive fracture strength of about 600 MPa compared to those of the as-cast specimen. Zhao et al. [19] conducted annealing treatments on Al_0.3CoCrFeNi HEA at 900 °C/30 min and 1050 °C/20 min. The results showed that the microstructure of the samples obtained from the former exhibited grain refinement, while the samples obtained from the latter exhibited obvious twinning and equiaxed grains. Moreover, the recovery and recrystallization during annealing resulted in the elimination of dislocations, leading to a decrease in the hardness of both samples. Stepanov et al. [20] successfully prepared a CoCrFeMnNiV_x high-entropy alloy system using vacuum arc furnace melting technology, and they studied the microstructure and mechanical properties of the alloys as-cast and after heat treatment. The results showed that without the addition of V, both the as-cast and annealed samples (1000 °C/24 h) exhibited an FCC structure. And, in the case of x = 0.25, trace amounts of σ phase began to precipitate after heat treatment. The above two alloys exhibited good ductility both as-cast and after heat treatment. But, when x > 0.5, precipitation (σ phase) also began to occur in the as-cast sample, and the σ phase became more abundant in the samples after heat treatment, resulting in a decrease in the elongation of the alloys but a significant increase in strength. However, research on the annealing state of the AlCoFeNiV high-entropy alloy series is still quite rare. Therefore, in this work, Al_0.6CoFeNi₂V_0.5 high-entropy alloy was successfully designed and prepared via a nonconsumable arc-melting process, and it was annealed at different temperatures (600 °C, 800 °C, or 1000 °C) for 4 h. The microstructure and mechanical properties of this alloy, both as-cast and annealed, were further studied.

2. Materials and Methods

A nonconsumable vacuum arc-melting furnace (DHL-1250, Shanghai Mengting Instrument Equipment Co., LTD, Shanghai, China) was used to prepare the Al_0.6CoFeNi₂V_0.5 high-entropy alloy system. Before melting, this HEA needed batching, and the alloy batching process was as follows: The raw Al, Co, Fe, Ni, and V materials were placed in alcohol and cleaned via ultrasonic vibration to remove surface impurities, which was followed by air-drying treatment. Afterward, based on the composition of each element in the HEA, an analytical electronic balance was used to match the designed high-entropy alloy, with an accuracy of 0.001 g. The nonconsumable arc-melting process was used to melt HEA, and the specific operation process was as follows: The obtained alloy materials were placed in a water-cooled Cu crucible, and the vacuum was continuously maintained at 3 × 10⁻³ Pa. Then, the pressure was adjusted to 0.05 MPa using recoil argon gas. To further remove residual oxygen in the vacuum chamber, the preset pure Ti block was melted more than twice to remove residual oxygen, and then the HEA was melted. After each melting, the obtained high-entropy alloy was thoroughly cooled, then flipped over, and further melted. To ensure the HEA had a uniform composition, it was melted at least 5 times to obtain a button-shaped alloy ingot with a uniform weight of 35 g, a diameter of about 25 mm, and a thickness of about 15 mm. Subsequently, a muffle furnace capable of heating up to 1200 °C was utilized to perform insulation air-cooled annealing at different temperatures of the obtained HEA ingots, with the aim of eliminating defects in the as-cast high-entropy alloys and observing the formation of new phases during the annealing process, investigating the impact of annealing temperature on the microstructure and macroscopic mechanical properties of the HEA. Additionally, the annealing time was 4 h, the annealing temperature was set to 600 °C, 800 °C, or 1000 °C, and the cooling was performed via air-cooling.

The detection of the crystal structure of the alloy in this study was carried out on an Empyrean X-ray diffractometer produced by Panaco in the Alemlo of Netherlands, which was combined with jade 6.5 software for analysis of the crystal structure. The main parameters during this experiment were an extreme voltage of 40 kV, a tube current of 40 mA, a maximum power of the copper target material of 2.2 kW, a scanning angle range from 20° to 100°, and a scanning step size of 0.0167. Before sample testing, it was necessary to perform surface pretreatment of the HEA samples. Sandpapers with different particle sizes were sequentially used to polish the high-entropy alloy samples, 0.5 μm diamond polishing paste was used for polishing treatment, and ultrasonic cleaning was used to clean the surface of the samples before testing in order to remove surface impurities and minimize the possible errors in this experimental process. In order to achieve more accurate morphology and composition measurements of the high-entropy alloy samples, electron probe microanalysis (EPMA) was used to characterize the microstructure of the HEA samples, using a JXA-8530F field-emission electron probe produced by JEOL in Tokyo of Japan. There was no need for corrosion treatment during the electron probe characterization of the high-entropy alloy samples, which effectively avoided the occurrence of false phases in the HEA samples during the corrosion process, thereby improving the accuracy and efficiency of the microstructure analysis.

An MH-5L Vickers hardness tester produced by Shanghai Yiheng Scientific Instrument Co., LTD in Shanghai of China was used to test the hardness of the HEA samples. The sample was pretreated before testing. The wire cutting method was used to cut as-cast HEA ingots into elliptical columns with parallel upper and lower surfaces, and sandpapers of different particle sizes were used to polish the upper and lower surfaces of the samples. Then, 1.5 μm diamond polishing paste was used to polish the HEA samples. Finally, the hardness of the samples was measured using a load of 1000 g and a loading time of 15 s. Each sample was tested seven times, and, after removing the maximum and minimum values, the average hardness value was taken as the final hardness value of the HEA samples. In addition, the HEA ingots were cut with wire to obtain Φ5 mm × 10 mm compression rods. Then, the surfaces of the compression rods were polished using sandpaper with different particle sizes (in the order of #80, #200, #600, #800), and the compression experiments were conducted on a UTM4000 electronic universal testing machine produced by Shandong Wanchen testing machine Co., LTD in Shandong of China. The maximum experimental force of the equipment was 100 kN, and the test for each sample was repeated three times. Finally, room-temperature compression stress–strain curves of the tested samples were drawn using Origin 8.0 software to obtain the mechanical performance parameters of the HEA samples.

3. Results and Discussion

3.1. Phase Constitution

Figure 1 shows the XRD patterns of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy, both as-cast and annealed at different temperatures for 4 h. It can be seen that the Al_0.6CoFeNi₂V_0.5 alloy had a dual-phase structure of FCC+BCC in the as-cast state. After annealing at 600 °C, it still had an FCC+BCC structure. By continuing to raise the temperature to 800 °C, after comparison with the standard PDF card library, it was found that the newly appeared peak positions in the XRD spectrum corresponded to the diffraction peak positions of Al₃V-type metal compounds. And, at 1000 °C, the Al_0.6CoFeNi₂V_0.5 alloy still had an FCC+BCC structure. The main reason for this was that the Al₃V metal compound dissolved at 1000 °C and could not precipitate in time under rapid air-cooling conditions. According to the literature [21], the formation of a solid solution in high-entropy alloy systems can be predicted by the enthalpy of mixing (ΔH_mix) and atomic radius difference (δ). Their specific calculations are shown in Equations (1) and (2).

$$\mathsf{\delta }=100\sqrt{{\sum }_{\mathrm{i}=1}^{\mathrm{n}}{{\mathrm{c}}_{\mathrm{i}}\left(1-\frac{{\mathrm{r}}_{\mathrm{i}}}{\mathrm{r}}\right)}^2}$$

(1)

$$∆H=4{\sum }_{\mathrm{i}=1,\mathrm{j}\ne 1}^{\mathrm{n}}{∆H}_{mix}^{i j}{\mathrm{c}}_{\mathrm{i}}{\mathrm{c}}_{\mathrm{j}}$$

(2)

In the above formulas, c_i is the molar percentage of element i, r_i is the atomic radius of element i, and r is the average atomic radius of all elements. ${∆H}_{mix}^{i j}$ is the mixing enthalpy of binary alloys. And, based on these, the parameters of the high-entropy Al_0.6CoFeNi₂V_0.5 alloy were δ = 4.87% and ΔH_mix = −9.04 kJ/mol. Research showed that when 0 < δ < 6.6%, −15 < ΔH_mix < 5 kJ/mol, stable solid solutions form [22]. The calculation results of the HEA system in this work were all within the above numerical range, so it could be concluded that solid solution phases can be formed in the Al_0.6CoFeNi₂V_0.5 alloy system. This is consistent with the XRD results mentioned above.

3.2. Microstructure

Figure 2 shows the scanning electron microscopy images of the as-cast Al_0.6CoFeNi₂V_0.5 high-entropy alloy and EPMA images of the annealed Al_0.6CoFeNi₂V_0.5 alloy. It can be seen that all HEA samples, both as-cast and annealed, had dendritic microstructures, and the main reasons for this phenomenon were as follows: Firstly, the Al_0.6CoFeNi₂V_0.5 high-entropy alloy contained a large number of elements; during the solidification process, the elements with higher melting points solidify first, while the elements with lower melting points accumulate at the front of the solidification interface, reducing the actual solidification temperature of the liquid phase at the front of the solidification interface and causing undercooling of the composition. In addition, due to the fact that the HEA samples were prepared in a water-cooled crucible, their cooling rates were fast; thus, they were in the nonequilibrium state. The growth mode was nonsmall plane growth [23]. During the process of crystal growth in nonsmall planes, the atoms gather at the solid–liquid interface with macroscopic smoothness but microscopic roughness, leading to the formation of columnar dendrites.

More specifically, the as-cast HEA sample had a very clear contrast between the light (A) and dark (B) regions, and the volume fraction of the former (83.51%) was much higher than that of the latter (16.49%), as shown in Figure 2a. As is well known, in XRD results, the higher the peak intensity of a phase, the higher its volume fraction [24]. In Figure 1, the intensity of the peak corresponding to the FCC phase is stronger than that of the BCC phase; therefore, the volume fraction of the former was much larger than that of the latter. Based on these results, it could preliminarily inferred that “region A”, which had a larger volume fraction, was the FCC phase, while “region B”, which had a smaller volume fraction, was the BCC phase. After further analysis of the alloy composition (

Table 1. Compositions of different regions (Figure 2) in Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples (at.%).

Sample	Region	Al	Co	Fe	Ni	V
As-cast	A	17.10	21.20	22.40	36.10	3.20
	B	25.50	14.00	14.70	38.20	7.60
600 °C	A	17.31	21.37	21.45	36.42	3.45
	B	26.13	14.30	14.21	38.13	7.23
800 °C	A	17.53	21.58	21.13	36.52	3.24
	B	25.19	15.33	13.26	39.03	7.19
	C	27.06	13.93	11.91	38.07	9.03
1000 °C	A	17.67	21.09	21.30	36.85	3.09
	B	25.68	14.66	14.35	38.23	7.08

), it was found that “region A” was enriched with Co and Fe. Gong et al. reported that a solid FCC solution is a Co–Fe-enriched phase [25]; this further confirms the rationality of the result that “region A” was the FCC phase. Additionally, Table 1 shows that “region B” was rich in Al and Ni elements; combining this with the XRD results, it could be inferred that “region B” was the BCC solid solution phase, which was also found by Ivanisenko et al. [26]. Additionally, according to the mixing enthalpy between the elements in

Table 2. The mixing enthalpy $∆H_{mix}$ of atom pairs in Al0.6CoFeNi2V0.5 high-entropy alloy (kJ/mol).

Element	Al	Co	Fe	Ni	V
Al	…	−10	−11	−22	−16
Co		…	−1	0	−14
Fe			…	−2	−7
Ni				…	−18
V					…

, it can be seen that Al and Ni have the most negative mixing enthalpy (−22 kJ/mol), so they easily enriched the B region and formed the BCC phase. Thus, “region A” was the FCC phase, while “region B” was the BCC phase in the as-cast HEA sample.

As shown in Figure 2b, after heat treatment at 600 °C, there was no significant change in the morphology compared to that of the as-cast state, and it still had a two-phase mixture of FCC+BCC, corresponding to the XRD pattern in Figure 1. In Figure 2c, it can be seen that there were some fine precipitates in the BCC phase. Figure 3 shows a local magnified electron probe image of the backscattering of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy after annealing at 800 °C, as well as surface scan images of various elements. From Figure 3, it can be seen that the precipitated needle-like particles were enriched with Al and V. Moreover, it can be further seen from Table 1 that the content ratio of Al and V was 3:1; thus, combining this result with the XRD results (Figure 1), it could be inferred that “region C (precipitate)” contained an Al₃V-type intermetallic compound. This was because the V mainly dissolved in the Al–Ni-rich BCC phase, which intensified the lattice distortion of the BCC phase; due to the relatively negative mixing enthalpy between V and Al (−16 kJ/mol, Table 2), the excess V atoms combined with Al to form an Al₃Ti-type intermetallic compound phase. After further calculation using Image-Pro Plus 6.0 software, it was found that the volume fraction of this needle-like precipitate was 7.37%. In Figure 2d and Figure 4, it can be seen that dispersed precipitates with the same contrast as BCC were precipitated in FCC. When vacuum arc-melting technology was used to prepare the Al_0.6CoFeNi₂V_0.5 high-entropy alloy, due to the rapid cooling effect at the bottom of the crystallizer, the undercooling at the interface was high, and various elements could not diffuse uniformly in time. The solidification process of the alloy deviated from the equilibrium conditions and was considered a no-equilibrium solidification process, so a metastable solid FCC solution was obtained; thus, during the high-temperature annealing process at 1000 °C, the solid BCC solution precipitated from the solid FCC solution. And, by using Image-Pro Plus 6.0 software for calculation, the volume fraction of the BCC phase was found to be as high as 21.81%. To confirm this substance more accurately, local magnification and elemental surface scanning analysis were performed on the Al_0.6CoFeNi₂V_0.5 high-entropy alloy annealed at 1000 °C using electron probes. The results are shown in Figure 4. It can be seen that the dispersed precipitates were completely identical in composition to the BCC phase, indicating that the alloy was annealed at 1000 °C, which promoted the formation of dispersed micro- and nano-scale BCC precipitates.

3.3. Mechanical Properties

Figure 5 shows the compressive stress–strain curves of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples, both as-cast and annealed. And, the rational performance parameters of the as-cast and annealed Al_0.6CoFeNi₂V_0.5 alloy samples are shown in

Table 3. Mechanical property parameters of as-cast and annealed Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples.

Sample	Yield Strength/MPa	Vickers Hardness/HV
As-cast	426.6	263.7
600 °C	353.4	216.9
800 °C	383.8	251.8
1000 °C	401.1	261.3

. The yield strength and hardness of as-cast alloy is 426.6 MPa and 263.7 HV. After annealing at 600 °C, 800 °C, and 1000 °C, the yield strengths of the Al_0.6CoFeNi₂V_0.5 high-entropy alloys decreased to 353.4 MPa, 383.8 MPa, and 401.1 MPa; the hardness was 216.9 HV, 251.8 HV, and 261.3 HV, respectively. It can be seen that compared with the as-cast alloy, the yield strength and hardness of the annealed samples were lower. This because the as cast HEA was prepared by rapid cooling in a water-cooled copper crucible after being melted in a vacuum arc furnace. There was significant internal stress present in this alloy, which was eliminated through annealing. More specifically, the annealing process can easily lead to an increase in grain size and a decrease in the number of grain boundaries, and the annealing process is accompanied by the movement of dislocations; thus, some dislocations would be expelled from the crystal, leading to a decrease in dislocation arrangements. Finally, the strength and hardness of the HEA in the as-cast state were higher than those in the annealed state. In addition, after annealing, the strength and hardness of the HEA samples were higher; the specific reasons for this are as follows: As is well known, the increase in the strength of high-entropy alloys is the result of multiple strengthening mechanisms, such as fine-grained strengthening (σ_G), dislocation strengthening (σ_D), dispersion strengthening (σ_P), and solid solution strengthening (σ_C). The yield strength ($\mathsf{\sigma }$_y) can be expressed as:

$${\mathsf{\sigma }}_{\mathrm{y}}={\mathsf{\sigma }}_0+{\mathsf{\sigma }}_{\mathrm{G}}+{\mathsf{\sigma }}_{\mathrm{D}}+{\mathsf{\sigma }}_{\mathrm{P}}+{\mathsf{\sigma }}_{\mathrm{C}}$$

(3)

In the above equation, σ₀ represents the intrinsic strength of alloys. The intrinsic strength of an alloy is mainly controlled by characteristics such as interatomic bonding energy; the other strengthening methods are determined by the microstructure of the alloys. As for the Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples, when the annealing temperature was 600 °C, it was composed of solid FCC and BCC solutions, mainly exhibiting the solid solution strengthening mechanism. When the annealing temperature was 800 °C, in addition to the solid solution strengthening effect of the two solid solutions mentioned above, there was also diffusion strengthening, caused by the precipitation of a new Al₃V-type phase with a volume fraction of 7.37%. When the annealing temperature was 1000 °C, a larger volume fraction of the BCC phase (21.81%) diffused and precipitated from the FCC phase, indicating that in addition to solid solution strengthening, the sample also had stronger dispersion strengthening. Therefore, with the increase in annealing temperature, the hardness and strength of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples showed a gradually increasing trend. Furthermore, it is worth noting that after annealing at 1000 °C for 4 h, the yield strength of the Al_0.6CoFeNi₂V_0.5 HEA decreased the least (about 5.9%), almost on par with that of the as-cast sample. This was mainly attributed to the dispersion strengthening effect caused by the dispersion precipitation of the BCC phase with a larger volume fraction (21.81%) from the FCC phase. Wang et al. [27] obtained similar results after comparing the mechanical properties of Al_0.4CoFeNiTi_0.6 high-entropy alloy before and after annealing treatment. That is, after annealing under the same conditions (1000 °C/4 h), the yield strength of this alloy was lower, but the magnitude of the decrease was much larger than that of the Al_0.6CoFeNi₂V_0.5 HEA in this study, which means that the mechanical properties of the latter are more stable before and after annealing. All these results promote the research and development of new HEA materials that bear compressive loads, such as columns in large factory buildings, supports for cranes, and clamping bolts for rolling mills in practical mechanical engineering.

4. Conclusions

In this study, a Al_0.6CoFeNi₂V_0.5 high-entropy alloy was successfully designed and prepared, and it was annealed at 600 °C, 800 °C, and 1000 °C for 4 h. The microstructure and mechanical properties were studied. The following conclusions were drawn:

• The as-cast Al_0.6CoFeNi₂V_0.5 high-entropy alloy had an FCC+BCC structure, and no phase transformation occurred during annealing at 600 °C. In addition, hard Al₃V-type metal compounds precipitated during annealing at 800 °C, and BCC particles precipitated in the FCC matrix during annealing at 1000 °C.
• Compared with the as-cast sample, the strength and hardness of the annealed Al_0.6CoFeNi₂V_0.5 high-entropy alloy samples were all lower after annealing because annealing eliminated the internal stress in this alloy.
• As the annealing temperature increased, the strength and hardness of the Al_0.6CoFeNi₂V_0.5 high-entropy alloy gradually increased. This was because new hard Al₃V-type intermetallic compound precipitates (the volume fraction was 7.37%) formed at the annealing temperature of 800 °C, which produced the “second phase strengthening” effect. The larger volume fraction of the hard and fine BCC phase (21.81%) diffusely precipitated at the annealing temperature of 1000 °C; meanwhile, the precipitation of this BCC phase produced a “second phase strengthening” effect, which also led to “solid solution strengthening”, ultimately resulting in enhanced hardness and strength.