Mechanical Alloying-RF Plasma Spheroidization for Preparing Spherical Powder of WMoTaNbV Refractory High-Entropy Alloy
Release time:
2026-01-29
High-entropy alloys (HEAs) typically contain five or more major elements, with each constituent present in equimolar or near-equimolar ratios. This alloy design concept was first proposed in 2004 by Professor Yue Junwei and his research team at National Tsing Hua University, Taiwan [1]. Unlike traditional alloys, HEAs exhibit high-entropy effects, hysteretic diffusion effects, lattice distortion effects, and “cocktail” effects, demonstrating outstanding mechanical properties, corrosion resistance, and thermal stability. Refractory high-entropy alloys (RHEAs) were introduced in 2010 by Professor Senkov from the United States, building upon the high-entropy alloy design philosophy [2]. Based on high-entropy alloys, RHEAs are alloy systems obtained by mixing high-melting-point elements in equimolar or near-equimolar ratios. Professor Senkov [3] first synthesized WMoTaNb and WMoTaNbV alloys via vacuum melting. RHEAs exhibit superior high-temperature strength, oxidation resistance, and phase stability. In recent years, RHEAs have achieved significant research breakthroughs in high-temperature, high-pressure applications across aerospace, energy, and nuclear industries. Their superior high-temperature stability, strength, and corrosion resistance position them as strong candidates to replace traditional high-temperature materials such as nickel-based alloys and molybdenum alloys.
Current RHEA fabrication techniques primarily include casting methods [3], powder metallurgy [4], and additive manufacturing processes [5-6]. The first two methods typically cannot directly form complex structural components. Additive manufacturing, by sequentially adding material layer by layer, enables personalized design and complex structural formation. It holds promise as a frontier technology for RHEA shaping, offering expanded possibilities for applications in high-temperature and extreme environments. Recent literature has also documented additive manufacturing of refractory high-entropy alloys. Kunce et al. [7] pioneered the use of laser near-net-shape forming to produce TiZrNbMoV refractory high-entropy alloy, revealing that the phase composition of TiZrNbMoV depends on the laser power applied during forming. Dobbelstein et al. [8] employed direct energy deposition technology using mixed element powders to fabricate the refractory high-entropy alloy MoNbTaW. Due to differing melting points among the refractory metals, a subsequent laser remelting step was required. They also observed that irregularly shaped particles exhibited poor powder focusing, a problem resolved by using spherical particles. Huber et al. [9] employed powder bed laser melting to in situ synthesize WMoTaNbV high-entropy alloys, using spherical particles for all elements except V powder. Although the final alloy exhibited a single-phase bcc structure and high density (99.8%), elemental segregation persisted within the samples. Currently, two approaches exist to prevent elemental segregation: one involves remelting the in-situ printed parts to achieve thorough elemental diffusion, while the other employs pre-alloyed powders for printing to avoid large-scale segregation. Lee et al. [10] pioneered the preparation of spherical WMoTaNbV pre-alloyed powders by vacuum melting WMoTaNbV alloy ingots, followed by hydrogenation-crushing-dehydrogenation treatment and plasma spheroidization to successfully produce spherical powders with a chemical composition of 20V-20Nb-20Mo-20Ta-20W (at%). Gu et al. [11] employed the aforementioned method to process spherical VNbMoTaW alloy powder for selective laser melting, producing bulk VNbMoTaW high-entropy alloys with superior mechanical properties compared to those fabricated via arc melting. Currently, commercially available RHEAs spherical powders for additive manufacturing are produced through this melting-crushing-sphericalization process. However, due to the lengthy preparation cycle and high costs, commercial powders are expensive. Therefore, it is necessary to explore new pathways for preparing refractory high-entropy alloy powders for additive manufacturing.
This work investigates the WMoTaNbV alloy system, proposing a combined approach of mechanical alloying and radiofrequency plasma spheroidization to produce low-impurity spherical WMoTaNbV refractory high-entropy alloy powder. First, five elemental powders were weighed in equal atomic ratios and ball-milled. The milled powders were then processed through a plasma spheroidization device to produce spherical WMoTaNbV refractory high-entropy alloy powder. The powder's microstructure, phase composition, morphology, particle size, and purity were examined. This study provides a reference route and fundamental data for the development and preparation of refractory high-entropy alloys.
1. Experiment
This experiment utilized irregular elemental powders as raw materials: W (particle size 74–180 µm, purity 99.9%), Mo (particle size 74–180 µm, purity 99.9%), Ta (particle size 74–180 µm, purity 99.9%), Nb (particle size 74–180 µm, purity 99.9%), and V (particle size 45–180 µm, purity 99.9%). The powder morphologies are shown in Figure 1.
Elemental raw material powders were weighed in equal atomic ratios and placed into ball milling jars. Both the jars and grinding media consisted of hard alloy balls. The ball mill (PLASMA-BM-L) employed equipment developed by Professor Zhu Min's team at South China University of Technology [12], operating similarly to a vibratory ball mill. The grinding parameters were as follows: Ball mill speed: 960 r/min; Ball-to-powder ratio: 50:1; Milling times set to 2, 3, 6, 10, and 14 h. Milling atmosphere: High-purity argon gas (99.99%).
The ball-milled powder was then fed into the radiofrequency plasma spheroidization system (TekSphero-40). The plasma used in spheroidization is a high-temperature plasma reaching 8000–10,000°C. As the powder passes through the plasma torch, it melts and undergoes surface tension-driven polycondensation into spheres, rapidly solidifying into spherical particles. Spheroidization process parameters are shown in Table 1.
Powder particle size was analyzed using a laser particle size analyzer (HORIBA LA960S). Powder phase composition was analyzed using an X-ray diffractometer (Philips-MPD X’Pert). Powder elemental content was analyzed using an X-ray fluorescence spectrometer (PANalytical Axios). powder microstructure was observed using scanning electron microscopy (Supra 40) and transmission electron microscopy (F200X); cross-sectional oxygen distribution was analyzed by electron probe microanalysis (EPMA); oxygen and carbon content in the powder was determined using a CS844 carbon-sulfur analyzer and an OHN836 oxygen-nitrogen-hydrogen analyzer.
2. Results and Discussion
2.1 Analysis of Mechanically Alloyed Powders
XRD patterns of the raw powder and powders prepared with different ball milling durations are shown in Figure 2. The figure indicates that characteristic diffraction peaks of each elemental powder are clearly detectable in the raw powder. After 2 h of ball milling, the Nb and Mo diffraction peaks gradually become difficult to distinguish, but the diffraction peak of the V element remains detectable. Consistent with literature [13], the Nb and Mo diffraction peaks did not disappear but broadened due to grain refinement and residual stresses during ball milling, causing overlap with the Ta and W peaks, respectively. As milling time increased, the diffraction peaks of the elemental components gradually weakened. At 10 h of ball milling, only Ta and W diffraction peaks were present; at 14 h, only a single set of diffraction peaks remained. Studies indicate [14-15] that during mechanical alloying, the alloying rate between elements correlates with melting point: lower melting points imply weaker atomic bonding and higher self-diffusion coefficients, thus accelerating alloying. This suggests all other elements solid-solve into the W lattice, forming a bcc solid solution structure. Since the atomic radius of W is smaller than those of Mo, Ta, and Nb, and W has the highest mass fraction in the powder, lattice distortion occurs in the W lattice upon solid solution formation. This increases the W lattice constant, causing the diffraction peaks in the XRD pattern to shift to the left relative to those of pure W. Calculations based on the XRD pattern indicate that the grain size of the bcc solid solution is 3.7 nm, with a lattice constant of 0.3187 nm.
Figures 3a–3e show the morphology of powders prepared at different ball milling times. As shown, at a ball milling time of 2 h, the powder undergoes severe deformation under the impact of grinding balls, resulting in flake-like particles. As milling time increases, cold welding occurs between these flake-like particles. After 6 h of milling, the powder transitions from flake-like to block-like morphology, with gradually increasing thickness. Following 14 h of milling, the powder morphology is predominantly block-like, with a significant reduction in particle size. Notably, some large particles persist. This is primarily due to work hardening from ball milling, which reduces plasticity and shrinks particle size, while cold welding causes small particles to recombine into larger ones. As shown in the particle size distribution (Figure 3f), increased milling time deepens fragmentation, with particles milled for 14 hours predominantly below 20 μm.
Figure 4 shows the TEM image of the refractory high-entropy alloy powder after 14 h of ball milling. Figure 4b clearly reveals nanoscale microcrystals. The selected area electron diffraction (SAED) pattern in this region exhibits a ring-like diffraction pattern, as shown in the inset of Figure 4b, indicating that the prepared powder has a polycrystalline structure with a body-centered cubic (bcc) phase structure. This structure is consistent with that determined from the XRD pattern shown in Figure 2. Based on measurements of interplanar spacings from the SAED pattern, the lattice parameter was calculated as 0.3203 nm, which agrees well with the lattice parameter of 0.3187 nm calculated from the XRD pattern. Figure 4c shows a dark-field TEM image of the ball-milled powder, clearly revealing numerous bright spots corresponding to nanoscale microcrystals. High-resolution TEM observation (Figure 4d) indicates that the microcrystalline grains marked with yellow ellipses have a size of approximately 5 nm and an interplanar spacing of 0.235 nm.
The ball-milling process may increase the oxygen and carbon content in the powder, which can significantly impair the performance of refractory high-entropy alloy components [16-17]. Figure 5 shows the variation in oxygen and carbon content in the powder with ball-milling time. Figure 5 indicates that the carbon content in the powder increases with extended ball milling time. This is attributed to carbon-containing milling debris generated during the process, which originates from the hard alloy milling media and tank and contaminates the powder, leading to elevated carbon levels. Regarding oxygen content, when milling time is less than 6 h, the oxygen content rises with increasing milling duration but subsequently stabilizes. This occurs because during the initial ball milling stage, powder fragmentation and refinement increase the specific surface area, leading to reactions with residual oxygen in the vessel and thus higher oxygen content. Subsequently, the powder particle size gradually stabilizes (as shown in Figure 3f), preventing further oxygen content increase. To ensure the performance of the refractory high-entropy alloy, the powder with the lowest oxygen and carbon content (ball milled for 2 hours) was selected for subsequent radiofrequency plasma spheroidization testing.
2.2 Analysis of Spheroidization in Ball-Milled Powders
Although the 2-hour ball-milled powder was not fully alloyed, the remelting process during RF plasma spheroidization promotes further alloying. XRD patterns before and after spheroidization are shown in Figure 6. As shown in Figure 6, the post-spheroidization powder exhibits a distinct bcc solid solution phase structure. However, characteristic peaks appear broadened, and minor impurity peaks are present. This phenomenon may result from partial powder failing to fully melt during passage through the RF plasma torch (as illustrated in Figure 8a). The incompletely melted powder likely consists of Ta with a higher melting point, as Ta powder achieves complete solid solution during the brief ball milling process. Although trace amounts of elemental impurities exist in the spheroidized powder, complete solid solution is achieved during 3D printing, resulting in a single bcc phase.
To further investigate the properties of the pre-spheroidized powder, SEM and EDS analyses were performed. Figure 7 shows the morphology and EDS elemental distribution of the powder after 2 hours of ball milling. Figure 7a reveals that the pre-spherical powder predominantly consists of flakes. EDS elemental analysis of the red dashed box in Figure 7b indicates uneven elemental distribution, with elemental contents shown in the table. Some elements still exist in elemental form.
Figure 8 shows the morphology of the powder after spheroidization and the EDS elemental distribution. As illustrated, the vast majority of the powder particles are spherical. During spheroidization, the powder undergoes four stages [18]: particle heat absorption, particle melting, liquid metal heat absorption and temperature rise, and powder evaporation due to heating. It is generally considered that the third stage represents the optimal condition for powder spheroidization [18]. Figure 8a reveals the presence of a small amount of flake-shaped powder. This occurs because a minor portion of the powder, influenced by factors such as carrier gas/sheath gas within the RF plasma, bypasses the plasma core, resulting in insufficient heating. As shown in Figure 8d, the elemental distribution within the spheroidized powder is uniform with no significant segregation, though the elemental content slightly deviates from the raw material composition.
While the previous analysis examined elemental distribution in individual particles, assessing overall powder composition uniformity is crucial for subsequent 3D printing consistency. Figure 9 shows EDS elemental mapping scans of multiple powders. The results indicate slight variations in elemental content between different powders, which correlates with the elemental composition of individual particles prior to spheroidization. To assess the overall elemental content of the spheroidized powder, quadrant sampling was employed for XRF analysis. The results, presented in Table 2, show that while the overall composition of the spheroidized powder deviates slightly from the raw material, this deviation is smaller than that observed in individual powders. This ensures the 3D printed parts maintain composition consistent with the high-entropy alloy definition.
To investigate oxygen distribution in the spheroidized powder, EPMA was further employed to analyze oxygen content distribution across cross-sections (Figure 10). The color transition from blue to orange indicates progressively increasing mass fractions. The results in Figure 10b confirm extremely low oxygen content within the powder particles. The prominent green areas near the cross-section result from the organic resin binder material, which contributes to the detected oxygen content. Furthermore, the XRD pattern shown in Figure 6 reveals no oxide diffraction peaks, further confirming the low oxygen content in the spheroidized powder.
The particle size distribution of the powder before and after spheronization is shown in Figure 11. As seen in Figure 11, the D₁₀ and D₅₀ particle sizes of the powder after spheronization decreased slightly, with D₅₀ = 55.9 µm, while the D₉₀ particle size decreased significantly, resulting in a markedly narrower powder size distribution. This may be attributed to the flake-like morphology of the powder after ball milling. When measuring non-spherical particles, laser particle size analyzers equate them to a combination of spherical particles of varying sizes, resulting in a measured particle size distribution that appears broader than the actual distribution. During the RF plasma spheroidization process, flake-like powder particles melt and coalesce into spheres. Assuming constant particle mass, the diameter is significantly reduced. Additionally, during spheroidization, smaller powder particles are prone to thermal vaporization and evaporation [19]. Melted fine particles may also adhere to the surfaces of larger particles.
Table 3 compares the physical properties of WMoTaNbV alloy powder before and after spheroidization. It can be observed that prior to spheroidization, the flake-like powder exhibits negligible flowability. Post-spheroidization, the Hall flow velocity reaches 8.4 s · (50 g)^(−1). Furthermore, both the bulk density and tapped density of the powder show substantial increases, with bulk density at 7.96 g · cm^(−3) and tapped density reaching 8.80 g · cm^(−3).
Figure 12 shows the oxygen and carbon content (mass fraction) of the powder before and after spheroidization. As shown in Figure 12, the oxygen and carbon content of the powder decreased significantly after spheroidization, to 0.054% and 0.021%, respectively. The reasons are: (1) The carbides formed during ball milling decomposed at high temperatures to form carbon, which reacted with the adsorbed oxygen on the powder and volatilized; (2) H₂ in the RF plasma sheath gas not only enhances plasma thermal conductivity [20] but also acts as a reducing agent [21] that reacts with oxides in the powder; (3) Small particles adhering to the powder surface vaporize and volatilize under the high temperature of the plasma.
The Gibbs free energy curves calculated using Outotec HSC Chemistry 9.5 software aid in understanding the reduction mechanism. Figure 13a shows the Gibbs free energy calculations for the oxidation of five elemental metals. It can be seen that a series of oxides including NbO, Ta₂O₅, NbO₂, Nb₂O₅, V₂O₅, WO₂, WO₃, MoO₂, and MoO₃. All oxidation reactions exhibit negative Gibbs free energy between room temperature and 3000°C, indicating spontaneous oxidation at these temperatures. However, as temperature increases beyond this point, the Gibbs free energy for all reactions except MoO₂ formation becomes positive, signifying that elevated temperatures reduce the driving force for oxidation. Generally, under identical process conditions, the more negative the Gibbs free energy of an equation, the more preferentially oxidation occurs, and the more stable the resulting oxide [22]. Therefore, after plasma treatment, oxygen primarily combines with Mo to form MoO₂. Figure 13b shows the Gibbs free energy diagram for carbide decomposition. It indicates that MoC can decompose at high temperatures to yield free carbon and elemental Mo, while other carbides exhibit resistance to decomposition at elevated temperatures, particularly TaC. Consequently, a certain amount of carbon remains present in the spheroidized powder. The Gibbs free energies for carbon and H₂ reacting with oxygen are shown in Figures 13c–13e. It can be observed that free carbon and H₂ can act as reducing agents to reduce oxides into metallic elements, H₂O, CO₂, and CO, thereby reducing the oxygen and carbon content in the powder. However, Figure 13e indicates that the hydrogen reduction of MoO₂ is difficult to occur in the high-temperature range. resulting in residual oxygen in the final spheroidized powder.
3. Conclusions
1) WMoTaNbV high-entropy alloy powder with a single bcc phase structure was successfully prepared via mechanical alloying. During the initial ball milling stage, the powder first underwent flattening as elemental atoms began diffusing. As milling time increased, cold welding and fragmentation progressively occurred, yielding a fully bcc-phase solid solution after 14 hours. Extended ball milling time correspondingly increased oxygen and carbon impurity content in the powder. To meet 3D printing requirements for impurity content and particle size, the optimal ball milling process for preparing pre-spherical powder was determined as: milling speed 960 r/min; ball-to-powder ratio 50:1; milling time 2 h.
2) Spherical WMoTaNbV refractory high-entropy alloy powder can be prepared via radiofrequency plasma spheroidization. Spherodization parameters selected were: power 40 kW; powder feed rate 20 g/min; reactor internal pressure 102.7 kPa; hydrogen flow rate of 5 L/min. The resulting powder exhibits high sphericity, smooth surface, and absence of satellite particles. The Hall flow rate reaches 8.4 s/(50 g)⁻¹, with a bulk density of 7.96 g·cm⁻³, tapped density of 8.80 g·cm⁻³, oxygen content of 0.054%, and carbon content of 0.021%. The excellent flowability, high tapped density, and low oxygen content of this spheroidized powder render it suitable for 3D printing applications.
Reference: Rare Metal Materials and Engineering, Vol. 53, No. 12; Mechanical Alloying-RF Plasma Spheroidization for Preparing Spherical Powders of Refractory High-Entropy Alloy WMoTaNbV; Wang Fanqiang 1,2, Shi Qi 2, Liu Xin 2, Liu Binbin 3, Tan Chong 2, Xie Huanwen 2, Shen Zhengyan 2, Zeng Meiqin 1
Mechanical Alloying - Radio Frequency Plasma Spheroidization for Preparing WMoTaNbV Refractory High-Entropy Alloy Spherical Powders is a core technology for producing high-end refractory powders tailored for additive manufacturing. Building upon our foundation in selling spherical powders of rare refractory metals, we offer complementary professional 3D printing services to create integrated solutions. This approach overcomes key technical challenges in powder production, enabling full process quality control from powder to formed parts. As demand for refractory high-entropy alloys continues to grow in advanced manufacturing sectors, we will deepen technological innovation. Through high-quality powder products and comprehensive forming services, we deliver superior material and processing solutions to support the advancement of high-end manufacturing industries. For further product details, please contact our Sales Manager, Sukie Zhu, at +86 13378626726.
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