Preparation of Bulk TiZrNbMoV and NbTiAlTaV High-Entropy Alloys by Powder Sintering

: The refractory HEAs block material was prepared by powder sintering, using an equal atomic proportion of mixed TiZrNbMoV and NbTiAlTaV metal powder raw materials. The phase was analyzed, using an XRD. The microstructure of the specimen was observed, employing a scanning electron microscope, and the compressive strength of the specimen was measured, using an electronic universal testing machine. The results showed that the bulk cubic alloy structure was obtained by sintering at 1300 ◦ C and 30 MPa for 4 h, and a small amount of complex metal compounds were contained. According to the pore distribution, the formed microstructure can be divided into dense and porous zones. At a compression rate of 10 − 4 s − 1 , the yield strengths of TiZrNbMoV and NbTiAlTaV alloys are 1201 and 700 MPa, respectively.


Introduction
High-entropy alloys (HEAs) are alloys containing several elements that violate the usual design constraints of one or two alloying components. They are distinguished by multi-component and multi-principal components that combine to produce nanoscale composite materials. [1] The lattice locations of multi-component HEAs are displaced to varying degrees due to the size difference of multi-principal element atoms, resulting in lattice distortion, which significantly enhances the microstructure and characteristics of the alloys [2][3][4].
In recent years, the study of HEAs has progressed to a new and interesting stage. The study emphasis of HEAs ranges from the crystal structures and mechanical characteristics of HEAs [5], such as hardness, wear resistance, tensile and compressive strengths, to magnetic properties, such as soft magnetic and hard magnetic behavior, magnetic refrigeration, and wave absorption, as well as HEAs film for photo-thermal conversion. Furthermore, the related HEAs preparation technology has been continually improved in response to meet the needs of alloys, and the preparation of eutectic HEAs, for hydrogen storage, high-entropy cemented carbides and coatings, refractory high-entropy alloys (RHEAs), irradiation-resistant HEAs, corrosion-resistant HEAs, and other special properties of HEAs [6][7][8]. The HEAs preparation and forming technology may also be used to develop a new classification system for HEAs [9], which can be separated into casting HEAs [10,11], thin-film HEAs [12][13][14][15][16], three-dimensionally (3D)-printed [17][18][19], deformation HEAs [20], powder-metallurgy HEAs [21,22], and numbered based on their phase type. Table 1 lists a few of the numbered alloys.
Vacuum arc melting is generally used to produce high entropy alloys. Although alloy casting has been industrialized, issues such as shrinkage cavity and segregation must still be addressed in the manufacturing and processing of alloys [23]. As a result, the study and development of material preparation technology is equally critical to promoting material application. Powder metallurgy has recently been used in the preparation of high entropy alloys [24].
The raw material powder is compacted by spark plasma sintering (SPS) and hotpressing sintering, and an alloy with uniform element distribution and no apparent segregation is produced [21,22]. The general SPS method is utilized in powder metallurgy to produce TaNbVTi series refractory HEAs. SPS with low temperature, short sintering time, and quick cooling speed may produce an alloy with fine grain and uniform structure. Controlling the temperature of SPS improves the strength and form of the alloy, but also limits the size of the produced alloy. Powder sintering can be used to produce the desired form and size of the alloy. By hot pressing and sintering gas atomized powders, Liu et al., [25] created densified FeCoCrNi high entropy alloy bars. The microstructure of the alloy reveals compact, fine, and homogeneous grains, a tensile strength of 750 MPa, an elongation of more than 50%, and mechanical characteristics superior to those of FeCoCrNi HEAs formed by conventional melting [25]. Powder metallurgy typically gets fine powder by mechanical alloying, and the chemical sequence of mechanical alloy is complicated and time-consuming, with contaminants injected into the process. The benefit of sintering is evident for fcc alloys and elements with low melting points, but it is difficult to produce powder by mechanical alloying and atomization for refractory element powder, such as Ti, Zr, and so on. As a result, after combining high purity element powder, we attempted to produce TiZrNbMoV and NbTiAlTaV HEAs by hot pressing sintering and evaluated the microstructure and mechanical characteristics.

Experimental Procedures
Multi-component TiZrNbMoV and NbTiAlTaV HEAs with equal atomic ratio were prepared by mixing high purity metal powder (the purity of 99.9%) with particle size of 10 microns (microstructure are depicted in Figure 1) homogeneously and hot pressing sintering at 30 MPa for 4 h at 1300 • C in high purity argon environment. For all the samples, densities of the sintered bulk alloys were measured by the Archimedes method [29]. The phase composition of the synthesized alloys was investigated by X-ray diffraction (XRD) (see Supplementary Materials) using a Ultima-IV 3 KW (Rigaku Corporation, Tokyo, Japan) diffractometer with Cu K α radiation. The microstructural investigation was with a Zeiss SUPRA a55 scanning electron microscope (SEM) (Carl Zeiss Microscopy GmbH, Jena, Germany) equipped with the energy-dispersive spectroscopy (EDS). Cylindrical specimens for compressive tests were 5.0 mm in diameter and 10.0 mm in height, and were tested with a WDW-200D electronic universal material testing machine at room temperature under strain rate of 10 −4 s −1 .   Figure 2 depicts the X-ray diffraction patterns of hot-pressing samples for the TiZrN-bMoV and NbTiAlTaV alloys series. The major diffraction peaks are determined to be a typical body centered cubic (BCC) phase. The peak contrast of the XRD diffraction pattern is represented by the positions of the respective peaks of the elements, Ti, Nb, Mo, V, and Ta, and includes several complicated compound peaks, such as AlTa, AlTi2, Mo9Ti4, Nb9Ti4, and Mo2Zr, but their intensities are low.  Figure 2 depicts the X-ray diffraction patterns of hot-pressing samples for the TiZrNbMoV and NbTiAlTaV alloys series. The major diffraction peaks are determined to be a typical body centered cubic (BCC) phase. The peak contrast of the XRD diffraction pattern is represented by the positions of the respective peaks of the elements, Ti, Nb, Mo, V, and Ta, and includes several complicated compound peaks, such as AlTa, AlTi 2 , Mo 9 Ti 4 , Nb 9 Ti 4 , and Mo 2 Zr, but their intensities are low.

Crystal Structure and Microstructure
Due to the small amount of metal compounds in the alloy, it cannot be fully detected in the XRD spectrum, according to the analysis. As a result, we used EDS and SEM to further investigate and observe the elemental distribution and microstructure morphology in the alloy. Figure 3 illustrates representative SEM images of the TiZrNbMoV and NbTiAlTaV alloys, which presents the distribution of ingredients and the phase composition of alloys The elements are not evenly distributed in these alloys, as can be seen, and there are dense and pore zones in the alloy, which would be described by Figure 4.
The results of backscatter indicated that different shades of color represent different element contents. As given in Figure  compounds, Point 2 is indeed mainly Nb, and the major element of point 3 is Ta. Point 4 is an AlV 4 intermetallic compound, and the composition is given in Table 2.  Due to the small amount of metal compounds in the alloy, it cannot be fully detected in the XRD spectrum, according to the analysis. As a result, we used EDS and SEM to further investigate and observe the elemental distribution and microstructure morphology in the alloy. Figure 3 illustrates representative SEM images of the TiZrNbMoV and NbTiAlTaV alloys, which presents the distribution of ingredients and the phase composition of alloys The elements are not evenly distributed in these alloys, as can be seen, and there are dense and pore zones in the alloy, which would be described by Figure 4.    The results of backscatter indicated that different shades of color represent differe element contents. As given in Figure 4a, the main component of Point 1 is the Ti6Zr compound, the main component of Point 2 would be Nb, and the main component Point 3 is V. Figure 4b describes that Point 1 is mainly composed of AlTi2 intermetal compounds, Point 2 is indeed mainly Nb, and the major element of point 3 is Ta. Poin is an AlV4 intermetallic compound, and the composition is given in Table 2.

Determination of Porosity
The Archimedes drainage method is used to measure the density of the samp shown in Equation (1). (1

Determination of Porosity
The Archimedes drainage method is used to measure the density of the sample, shown in Equation (1).
where m is the mass of the sample in the air; m a is the mass measured by the Archimedes drainage method [29]; ρ is the density measured by the Archimedes drainage method; and ρ l is the density of alcohol, which is 0.789 g/cm 3 . Due to the existence of pores in the material structure, alcohol was used in the drainage method for density measurements, which is easy to volatilize, and the sample drying is more convenient to reduce the measurement error. Before that, the theoretical densities of the samples were calculated as follows: 7.1660 g/cm 3 and 7.8826 g/cm 3 , and then the porosity is calculated. The calculation method of porosity is shown in Equation (2).
where η (%) is the porosity of the sintered HEAs, ρ (g/cm 3 ) is the actual density of the measured sample, and ρ th (g/cm 3 ) is the theoretical density of the sintered HEAs. The data measured by the Archimedes drainage method are shown in Table 3.  Figure 5 illustrates the compressive engineering stress-strain curve for alloys under a strain rate 10 −4 s −1 at room temperature. Among them, the TiZrNbMoV alloy exhibits a higher yield strength reaching 1201 MPa. In addition, the yield strength of the NbTiAlTaV alloy is 700 MPa. The elastic modulus of these two alloys are 27 and 23 GPa, respectively. Judging from the morphology of the compression fracture SEM images, which are shown in the Figure 6, the compression fracture is a typical brittle fracture with "river-like" or "stair-step" features. This trend is related to the crystal structures of the two alloys. The crystal structures of the two alloys are typical bcc structures. In the bcc crystal structure, the critical stress of the actuating slip system is larger and the number of slip systems is less than FCC, thus it shows the performance characteristics of high strength and low plasticity. Figure 5 illustrates the compressive engineering stress-strain curve for alloys under a strain rate 10 −4 s −1 at room temperature. Among them, the TiZrNbMoV alloy exhibits a higher yield strength reaching 1201 MPa. In addition, the yield strength of the NbTiAlTaV alloy is 700 MPa. The elastic modulus of these two alloys are 27 and 23 GPa, respectively. Judging from the morphology of the compression fracture SEM images, which are shown in the Figure 6, the compression fracture is a typical brittle fracture with "river-like" or "stair-step" features. This trend is related to the crystal structures of the two alloys. The crystal structures of the two alloys are typical bcc structures. In the bcc crystal structure, the critical stress of the actuating slip system is larger and the number of slip systems is less than FCC, thus it shows the performance characteristics of high strength and low plasticity.  The strength of the alloy prepared by hot-pressing sintering is significantly lower when compared to as-cast alloys [30,31]. The analysis of the strength of the hot-pressing sintering state alloy is the main reason for the due to the low preparation temperature of alloy, alloy body internal has not fully alloyed, causing the bonding force between the atoms that make up the elements to be smaller. Furthermore, the scanned fracture surface morphology images in Figure 6 shows that there are many macroscopic pores in the alloy, which are primarily formed by the expansion of air during the sintering process. Figure 7 depicts a schematic representation of the hot-pressing process and the prepared alloy to help you better understand the hot-pressing process. The computer creates the model of Figure 7c using 3D drawing software, so that the alloy produced by hot pressing may be seen more clearly. The real prepared sample is shown in Figure 7d; however, the form of the test sample was supported due to the test need.

Compressive Properties
morphology images in Figure 6 shows that there are many macroscopic pores in the alloy, which are primarily formed by the expansion of air during the sintering process. Figure 7 depicts a schematic representation of the hot-pressing process and the prepared alloy to help you better understand the hot-pressing process. The computer creates the model of Figure 7c using 3D drawing software, so that the alloy produced by hot pressing may be seen more clearly. The real prepared sample is shown in Figure 7d; however, the form of the test sample was supported due to the test need.

Discussion
In the experimental results, we give the alloys. The phases of the two alloys are mainly typical BCC and a small amount of other compound phases. There are porous and compact areas in the surrounding structure of the alloy. The yield strengths of the two alloys are 1201 and 700 MPa, respectively, and the elastic modulus are 27 and 23 GPa, respectively. These findings contradict assertions that powder metallurgy may generate dense and homogenous alloys. The majority of the refractory elements Ta, Nb, and Mo are agglomerated in the form of elemental in the alloy sintered at 1300 °C. Ti occurs as a body-centered cubic β-Ti phase and as a composite phase with Zr, Al. The goal of sintering metal powder is to produce a solid solution and compound via diffusion. Because the majority of the component elements are refractory elements with high melting points, and melting point of Al is relatively low, and the diffusion ability and rate increase with increasing sintering temperature, Ti-Al, Ta-Al compounds are produced preferentially. The property of Zr, V, on the other hand, is comparable to that of Ti, and a little quantity of solid solution is produced. It is believed that the solid solubility of Ti-Zr-V is low when sintered at 1300 °C, and when cooled in vacuum with the furnace, Zr and V are β stable phases, Ti progressively transforms to α -Ti in the cooling process, and the solid solubility in Ti declines, resulting in a very few solid solution phases. Nevertheless, due to inadequate sintering temperature, Nb, Mo, and Ta tends to produce elemental agglomeration. The microstructure also shows that the majority of the dense regions are composed of compounds and solid solutions generated by elements with low melting points and good diffusion capacity, whereas the holes are composed of elements with high melting points.
An alloy fracture may be defined as the process of crack development, propagation, and fracture. The fracture type of the two alloy groups in the experiment is the same, which belongs to cleavage brittle fracture. The pore region at the surface of the alloy forms stress concentration and microcracks when subjected to external force, and there is no strength response at the beginning of the stress-strain curve. When the load increases, the

Discussion
In the experimental results, we give the alloys. The phases of the two alloys are mainly typical BCC and a small amount of other compound phases. There are porous and compact areas in the surrounding structure of the alloy. The yield strengths of the two alloys are 1201 and 700 MPa, respectively, and the elastic modulus are 27 and 23 GPa, respectively. These findings contradict assertions that powder metallurgy may generate dense and homogenous alloys. The majority of the refractory elements Ta, Nb, and Mo are agglomerated in the form of elemental in the alloy sintered at 1300 • C. Ti occurs as a bodycentered cubic β-Ti phase and as a composite phase with Zr, Al. The goal of sintering metal powder is to produce a solid solution and compound via diffusion. Because the majority of the component elements are refractory elements with high melting points, and melting point of Al is relatively low, and the diffusion ability and rate increase with increasing sintering temperature, Ti-Al, Ta-Al compounds are produced preferentially. The property of Zr, V, on the other hand, is comparable to that of Ti, and a little quantity of solid solution is produced. It is believed that the solid solubility of Ti-Zr-V is low when sintered at 1300 • C, and when cooled in vacuum with the furnace, Zr and V are β stable phases, Ti progressively transforms to α -Ti in the cooling process, and the solid solubility in Ti declines, resulting in a very few solid solution phases. Nevertheless, due to inadequate sintering temperature, Nb, Mo, and Ta tends to produce elemental agglomeration. The microstructure also shows that the majority of the dense regions are composed of compounds and solid solutions generated by elements with low melting points and good diffusion capacity, whereas the holes are composed of elements with high melting points.
An alloy fracture may be defined as the process of crack development, propagation, and fracture. The fracture type of the two alloy groups in the experiment is the same, which belongs to cleavage brittle fracture. The pore region at the surface of the alloy forms stress concentration and microcracks when subjected to external force, and there is no strength response at the beginning of the stress-strain curve. When the load increases, the stress concentration near the pore tip becomes easier, promoting the crack's quick propagation. Because pores and microcracks already exist throughout the material, when subjected to external stress, the generated microcracks and pores rapidly grow and link from the surface to the interior, resulting in material fracture. As a result, pores and fractures frequently become the fracture source of stress concentration in the pore area, causing the material to fracture at lower stress and reducing strength, plasticity, and toughness.
Our results differ from those achieved with dense, high-strength alloys that are hotpressed and sintered after the elements are mixed directly with us without prior mechanical alloying. Therefore, the density and mechanical properties are moderate. However, our findings shed some light on the synthesis of sintered high entropy alloys. We discovered that the elastic modulus of the two alloys is similar to that of human bone, and TiZrNb has good biocompatibility [32]. Matching with bone modulus is a required requirement, and too much strength is present, which will prevent explants from adapting in vivo. Therefore, the production of porous materials by hot pressing and sintering element mixture powder opens the door to biomedical alloys. In addition, compared with mechanical alloying and atomized powder metallurgy, hot pressing sintering saves mechanical alloying time, does not introduce impurities in the process of hot pressing sintering, and avoids the danger of atomized powder production.

Conclusions
In this article, HEAs of TiZrNbMoV and NbTiAlTaV were created by powder hotpressing sintering. Structures of bcc phases were discovered in these alloys. Furthermore, alloys had a certain degree of strength, with porosity of 4.44% and 5.54%, respectively. During the sintering process, areas with varying densities and components were produced due to the various spreading abilities of the elements. More research is needed to explore the friction performance and applicability in biomedical materials and selflubricating materials.
Supplementary Materials: The following are available online at https://www.mdpi.com/article/ 10.3390/met11111748/s1, Figure S1: Stress-strain curve of AlxTiNbMoV and TiZrNbMoVx alloys (a) and variation of alloy plasticity with V atom ratio (b) [29,30].; Figure S2: XRD patterns of elements labeled with diffraction peaks.; Figure S3: Details of X-ray diffraction spectra of Nb elements.; Figure S4: Details of X-ray diffraction spectra of Ti elements.; Figure S5: Details of X-ray diffraction spectra of Ta elements.; Figure S6: Details of X-ray diffraction spectra of Mo elements.
Author Contributions: Writing-original draft preparation, Y.W.; supervision, Y.Z.; writing-review and editing, P.K.L. All authors have read and agreed to the published version of the manuscript.

Conflicts of Interest:
The author declares no conflict of interest.