Synthesis of high-entropy boride powders via boro/carbothermal reduction method

ABSTRACT In this paper, high-purity and fine (W0.2V0.2Ta0.2Nb0.2Ti0.2)B2 (HEB) powders were obtained via a simple high-energy ball milling-assisted boro/carbothermal reduction method at 1600°C. Results showed that the heating temperature and B source contents played a key role in the synthesis of purity of single-phase HEB powders. When excess 10 wt.% B4C was introduced, the single-phase HEB powders were synthesized at 1600°C, and its grain size was ~150 nm and it showed the hexagonal crystal structure of metal diborides. EDS results showed the high compositional uniformity of W, V, Ta, Nb, Ti and B elements in HEB powders. Through discussing the thermodynamics process related to these possible chemical reactions and combining the XRD results at different temperatures, it was obtained that higher temperature and enough B source were required to stimulate the formation of high-entropy phase. This work will be a vital step for the commercialization of high-entropy boride powders.


Introduction
The high-entropy borides, as a kind of high-entropy ceramic families, have received wide attention for their potential application in many fields such as cutting tool, metallurgy, thermal protection coating, ultrahigh temperature materials and nuclear reactors, etc. This is mainly attributed to their excellent physical and chemical properties including high hardness, chemical inertness and superior electrochemical properties, etc. [1][2][3][4][5].
In recent years, many methods have been successfully used to prepared the high-entropy borides, including the direct synthesis by high-temperature treatment using metal diboride powders as raw materials [1,6], the borothermal reduction method [7] and boro/carbothermal reduction method from oxides materials [8][9][10][11][12][13][14], and self-propagating hightemperature synthesis from metal and B powders [15], molten salt synthesis method [16] and so on. Among these methods, the boro/carbothermal reduction method can realize the large-scale production of high entropy diboride powders, which has been verified in preparing single phase diboride powders [17,18]. Liu et al. [11] used this method to successfully prepare high entropy diboride (Hf 0.25 Ta 0. 25 Nb 0. 25 Ti 0.25 )B 2 powders. However, lower temperature below than 2073 K leaded to the incompletion of boro/carbothermal reduction process, not forming the single phase high-entropy boride powders. High temperature of 2073 K was needed to realize the pure high-entropy boride powders on the condition of excess 40 wt% B 4 C in raw powders, and the average particle size was 260 nm. As is known that high-temperature condition easily causes the formation of coarse grains, which goes against the densification and the improvement of properties in sintering process of high-entropy powders [19]. Therefore, the preparation of high-entropy boride powders with fine grain size at lower temperature is crucial. High-energy ball milling process can realize the powders to the nanoscale in powder processing, which offers the opportunities to increase the reaction kinetics of powders [20,21]. This refinement process can realize the synthesis of fine high-entropy boride powders, which has been successfully used to synthesize fine ZrB 2 powders [22,23]. Nowadays, the research about the preparation of high-entropy boride powders via a simple high-energy ball milling processassisted boro/carbothermal reduction is still rarely. Meanwhile, owing to the big compositional design space brought by the complexity and multiformity of transition metal elements, the researches about the synthesis of new kind of high-entropy diboride ceramics are still insufficient and need to further be extended.
Therefore, in this paper, (W 0.2 V 0.2 Ta 0.2 Nb 0.2 Ti 0.2 )B 2 (HEB) powders were successfully prepared by the highenergy ball milling (HEBM) assisted boro/carbothermal reduction method. The thermodynamic behavior of HEB precursor was analyzed by the thermodynamic calculations and TG-DSC process. The effects of the heating temperature and B 4 C content on the formation of as-prepared HEB powders were analyzed. The morphology and structure of as-prepared HEB powders were studied. Furthermore, the formation mechanism of as-prepared HEB powders was discussed. The purpose of this study was to provide a vital step for the commercialization of high-entropy boride powders.

Characterization
The thermal behavior of HEB precursor was studied by TG-DSC equipment (STA449F3, NETZSCH-Gerätebau GmbH, German). The phase composition of HEB precursor and as-prepared HEB powders was characterized using X-ray diffraction (XRD, D/max-2550VB+/PC, Japan). The morphologies and structure of HEB precursor and as-prepared HEB powders were analyzed by a scanning electron microscope (FESEM, S-4700, Hitachi, Tokyo, Japan) and a transition electron microscope (TEM, Tecnai F-20, FEI, Hillsboro, OR) with energy-dispersive X-ray analyzer system. The particle size distributions of HEB precursor and as-prepared HEB powders were statistically measured from about 200 grains.

Characterization of HEB precursor
XRD pattern and SEM image of HEB precursor powders prepared by HEBM are shown in Figure 1. Five metal oxides, B 4 C and C phase, are observed in Figure 1(a). Due to the HEBM process, the particles of the original powder appear nano-sized, which will cause the diffraction peaks to be significantly broadened. To confirm this speculation, SEM image and the particle size distribution of the as-prepared HEB precursor powders after HEBM are given in Figure 1(b). It is found that the HEB precursor powders are spherical, with fine and good dispersibility. The particle size of HEB precursor powders is mainly concentrated at 100-300 nm, and its D 50 is 199 nm, which is consistent with the SEM results. It is worth noting that about 1um of soft agglomerates are also observed in the SEM results. Figure 2 shows the TG-DTG-DSC curves of prepared HEB precursor powders. It is evidently seen that the TG curve almost shows no weight loss before 1100°C, meanwhile, no strong endothermic or exothermic peaks can be observed in DSC curve, which indicates that no reduction reaction has occurred. A slight weight loss is observed above 1100°C, as the release of gaseous product. As the temperature increases, the rate of weightlessness increases, as shown by the DTG Figure 1. XRD patterns (a); SEM image and particle size distribution (b) of HEB precursor.

Preparation of HEB powders
curve, reaching a maximum at 1400°C. Additionally, a strong endothermic peak at 1400°C can be evidently observed in the DSC curve, attributing to the boro/ carbothermal reaction of oxides. The flat TG curve indicates the almost completeness of boro/carbothermal reaction process of precursor. It is obtained that the ceramic yield is 54% at 1500°C. Moreover, a huge endothermic range exists at around 1400°C, which is believed to be the solid solution process of the obtained powders. Figure 3 shows the phase evolution of HEB precursor at different temperatures for 2 h. When the mixed powders are prepared based on the stoichiometric ratio, the XRD of obtained HEB powder is all composed of a dominant (W 0.2 V 0.2 Ta 0.2 Nb 0.2 Ti 0.2 )B 2 (AlB 2 prototype, space group P6/mmm, no. 191) phase and minor (W x , M 1-x )B (CrB prototype, space group Cmcm, no. 63) solid solution phase, meaning that the solid solution process is incomplete at 1500°C and 1600°C [24][25][26][27][28]. On the one hand, the formation of (W x , M 1-x )B solid solution phase may be caused by insufficient B source under the stoichiometric ratio of the raw materials [29,30]. Usually, a small amount of B 2 O 3 is found on the surface of B 4 C, which has a higher vapor pressure at high temperature. Additionally, influence of O 2 impurity in Argon as well as diffused through the corundum furnace tube on the amount of B 4 C cannot be ignored [13]. On the other hand, Qin [24] et al. found the dissolution and stability of soft WB 2 phase in high-entropy ceramics, the formation of (W x , M 1-x )B solid solution phase could be explained via the formation enthalpies between AlB 2 -type and CrB-type structures. The formation enthalpies (H N NB À H N NB2 ) are −0.277, −0.07 and −0.163 in eV/atom, for N = W, Nb and Ta, from the Material Project Database [31]. It can explain why the monoboride phase observed in HEB is (W x , M 1-x )B. Finally, although having the similar XRD patterns at 1500°C to 1600°C, combined with the DSC curve results, it is more appropriate to choose 1600°C to obtain the solid solution phase owing to that higher temperature offers bigger reaction energy for the formation of single-phase solid solution.
To avoid the formation of (W x , M 1-x )B solid solution phase, the HEB powders are further prepared at 1600°C by using excess amount of B 4 C. The XRD results of obtained HEB powders are shown in Figure 4. Obviously, with increase of B 4 C amount, the diffraction peaks of (W x , M 1-x )B solid solution phase gradually decreases. (W x , M 1-x )B solid solution phase disappears when the excess amount of B 4 C is 10 wt.%, indicating that excess amount of B 4 C accelerates the conversion of (W x , M 1-x )B solid solution. Furthermore, Table 1 lists the lattice parameters of single phase diboride and HEB. It can be clearly seen that the lattice parameters of HEB are almost equal to the average value of five individual borides, which further proves that the asobtained powders are (W 0.2 V 0.2 Ta 0.2 Nb 0.2 Ti 0.2 )B 2 .     Figure 5 shows the SEM image and particle size distribution of HEB powders prepared with excess 10 wt% B 4 C at 1600°C for 2 h. Figure 5(a) shows that obtained products has sphere-like particles and the particles are uniform and fine. Furthermore, the particle size distribution is shown in Figure 5(b), and the D 50 is about 151 nm. By comparison, it is found that the particle size is reduced from 199 nm of HEB precursor to 151 nm of the as-obtained HEB powders, indicating that the boro/carbothermal reduction decreases can reduce the particle size of HEB powders. Additionally, it is seen that some weak sintering necks have been formed between the high entropy diboride particles, indicating that the as-obtained products have higher sintering activity.
A representative TEM morphology of as-synthesized HEB powders with excess 10 wt.% B 4 C content synthesized at 1600°C is shown in Figure 6(a). The HEB powders show near-spherical shape and their average particle size (measured by 50 particles sizes) is ~151 nm, which has been confirmed in SEM image. Meanwhile, an amorphous layer (B 2 O 3 ) with ~2 nm is observed on the HEB powder surface. The EDS compositional maps of as-obtained HEB powders are also given. Ta, Nb, Ti, W and V elements are shown that the highly uniform distribution without element agglomeration is scanned in as-obtained HEB powders, further confirming that as-synthesized products are composed of (W 0.2 V 0.2 Ta 0.2 Nb 0.2 Ti 0.2 )B 2 .

Formation mechanism of HEB powders
The thermodynamic calculations are used to analyze the synthesis possibility of HEB powders. The Gibbs free energy change (ΔG θ R;T ) of the possible chemical reactions is conducted by using HSC 6.0 (a software to calculate the Gibbs free energy change, produced by the Metso Outotec Finland Oy). According to the raw materials, the possible reactions and the relationship between ΔG θ R;T and temperature are as follows: where ΔG M R;T actually stands for mixing Gibbs free energy change of HEB powder. Assuming that HEB powder is the ideal Raoultian solution, ΔG M R;T can be calculated by the Equation (8).
Where ΔS min actually stands for mixing entropy of HEB powders, which is expressed based on the following equation [3,14]: Where ideal gas constant R is 8.314, metal element species N is 5, and molar fraction of the ith metal element in sublattice is x i . From Equations (8) and (9), ΔG M R;T is calculated to be −4.5 T. Therefore, the mathematically expressed of ΔG θ R;T from reaction (7) can be shown as: The system may maintain a standard pressure due to the flowing Argon gas. The unknown reaction rate of chemical reactions makes it difficult to precisely evaluate or measure the partial pressure of CO gaseous product in the system [11]. As stated in other reports, assume that the partial pressure of CO is equal to the standard pressure [32]. Figure 7 shows the ΔG θ R;T of those possible reactions. It can be observed that ΔG θ R;T of all reactions is negative (T > 1400 K) indicating that they all proceed spontaneously. However, it can be seen that due to its very negative ΔG θ R;T , compared to reaction 6, the reaction 7 is prone to generate HEB. In short, it is possible to obtain HEB by reaction 6 from the thermodynamic aspect.
To clearly understand the formation process of HEB powders, Figure 8 shows the phase evolution of excess 10 wt.% B 4 C precursor at different temperatures for 2 h. At 1200°C, unreacted oxides, newly formed borides and oxide solid solutions are observed. In addition, obvious diffraction peaks of TMC (the transition metal carbide, TM = W, V, Ta, Ti, Nb) phase can be observed in samples obtained at 1200-1400°C. This indicated that in addition to reactions (1)(2)(3)(4)(5)(6)(7), the other reactions may also occur [11].
where TMO are the five oxides of transition metals. Obviously, compared with the reaction (1-5), Equation (11) can consume much more amount of B 4 C. What is more, the by-product B 2 O 3 has a high vapor pressure and thereby it will rapidly evaporate with flowing Ar gas [32,33], which can further accelerate the occurrence Equation (11) to consume much more amount of B 4 C. This explained the generation of TMC phase by the Equation (12). Compared with the diffraction peaks at 1300°C, diffraction peaks of TMC phase are significantly weakened at 1400°C, accompanied by the formation of TMB, indicating that the solid solution phase of borides in the sample begins to form. The wide diffraction peaks indicate that the solid solution process has not been completed at 1400°C, which is the superposition of multiple solid solution phase diffraction peaks (such as Ta-based and Nb-based solid solutions). In addition, in DSC curve, a strong endothermic peak observed at 1400°C means the formation of single-phase boride and main high entropy solid solution phase. When temperature is increased to 1500°C, HEB is the main phase, but traces of impurity peaks can also be observed, indicating that higher temperatures are required to stimulate the formation of high-entropy phase. The increasing temperature causes the transformation from solid solution of TMC to high-entropy diboride by accelerating the boron/carbothermal reduction reaction. The solid solution situation of the powders with the temperature is analyzed using unit cell parameters. The lattice parameter of the highentropy phase is obtained by adding the unit cell parameter values of five single-phase borides and taking the average value. This value can be approximately regarded as the theoretical value of the high-entropy phase. This method has been adopted by many researchers, such as Liu et al. [3] and Feng et al. [13], contributing to analyze the solid solution process. Finally, JADE is used to calculate the unit cell parameters of the products obtained at different temperatures. The data are shown in Table 2. According to the calculation results, the unit cell parameter a of the high-entropy phase gradually decreases with increasing temperature, and the value of c becomes larger due to the solid solution of W at 1600°C. At 1600 degrees, the values of a and c are close to the theoretically calculated values, further confirming pure HEB can be obtained at 1600°C.

Conclusion
Single-phase and fine high-entropy boride (HEB) powders were obtained by a simple and easy HEBM assisted boro/carbothermal reduction method. Results showed that the temperature and B source contents played a key role in the synthesis of purity HEB powders. When excess 10 wt.% B 4 C was introduced, the single-phase HEB powders were synthesized at 1600°C, and it showed the hexagonal crystal structure of metal diborides. The particle size was reduced from 199 nm of HEB precursor to 151 nm of the as-obtained HEB powders by the boro/carbothermal reduction. Meanwhile, the HEB powder had a compositional uniformity of W, V, Ta, Ti, Nb and B elements, which was confirmed by EDS analysis using the TEM. Through discussing the thermodynamics process related to these possible chemical reactions and combining the XRD results at different temperatures, it was obtained that higher temperature and enough B source were required to stimulate the formation of high-entropy phase. This work will be a vital step for the commercialization of high-entropy boride powders.