Understanding the phase evolution and elemental distribution in MoWTaNbVTix manufactured via powder metallurgical approach

doi:10.21203/rs.3.rs-4438323/v1

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Understanding the phase evolution and elemental distribution in MoWTaNbVTix manufactured via powder metallurgical approach

https://doi.org/10.21203/rs.3.rs-4438323/v1

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The high melting point and thermal stability of MoWTaNbVTi_x refractory high entropy alloys have prompted their potential for harsh temperature structure application. A low-cost manufacturing strategy is key to realizing this promise. Powder metallurgy is known to be a cost-efficient manufacturing route for manufacturing refractory alloys. Elemental distribution and phases that form during manufacturing are crucial to realize the desired properties. However, the phase evolution and elemental distribution in MoWTaNbVTi_x refractory high entropy alloys manufactured via powder metallurgy (PM) have not been rationalized earlier. The present study focuses on the phase evolution and elemental distribution analysis in MoWTaNbVTi_x manufactured by different PM approaches, pressureless sintering, hot isostatic pressing, and spark plasma sintering. A series of MoWTaNbVTi_x were manufactured using different PM manufacturing methods. Using intensive X-ray diffraction, microstructure, and elemental analysis, we reveal the formation of an FCC (TiTaNb)_0.53C_0.47 carbide and TiO_x oxide phase in the nonequiatomic BCC MoWTaNbV matrix in MoWTaNbVTi_x. Interestingly, the liquid phase sintering yields a BCC_LSS matrix with a lattice parameter of 3.176 Å, differently than the solid state sintered BCC_SSS matrix with a lattice parameter of 3.193 Å. The fraction of (TiTaNb)_0.53C_0.47 and TiO_x phases varies with Ti composition and PM processing routes. Nickel, as the sintering activator, enhances Ti diffusion in the matrix. Also, the hardness of the MoWTaNbVTi_x varies nonlinearly in a range of ~ 602- ~911 Hv with the ratio of oxide and carbide fraction. The hardness obtained in the PM approach is significantly higher than that observed for cast MoWTaNbVTi_x refractory high entropy alloys, reported in the literature.

Spark plasma sintering

Hot isostatic pressing

Pressureless sintering

Precipitation

Phase evolution

Elemental distribution

Refractory high entropy alloys

Powder metallurgy

Due to their excellent temperature capability, the refractory metal and alloy family, tungsten (W), molybdenum (Mo), niobium (Nb), tantalum (Ta), vanadium (V), have been potentially considered for the extreme range of temperature applications including jet engines[1], nuclear reactors[2, 3], space power and propulsion[1, 4]. However, the inherent limitation of manufacturability (difficulty in converting unfinished products to finished/semi-finished products) and high-temperature oxidation resistance limit their use in industrial applications[5–7]. Individual elements or binary/ternary of these elements provide limited advantages for many applications. Since Yeh et al. [8] have reported that a solid solution phase can form when multiple principal elements in the equimolar composition are alloyed, a new vast materials domain is open for exploring. The new domain of materials space is referred to as high entropy alloys (HEAs), as their configurational entropy is higher than their conventional counterparts. Due to this unusual mixing of elements, effects like high entropy, lattice distortion (due to atomic size mismatch), sluggish diffusion (due to lattice strain), and cocktail effects are expected in these material systems[9]. These provide new opportunities to potentially overcome many challenges of conventional refractory alloys. It opens vast space for exploring new alloys that mitigate the fundamental barriers of refractory alloys for high-temperature applications. HEAs, specifically those made from refractory elements, referred to as refractory high entropy alloys (RHEAs), are considered potential alternatives as high-temperature materials[10–13].

Senkov et al. [14] reported the first RHEAs based on equiatomic elements: Nb, Mo, Ta, W and Nb, Mo, Ta, W, V. The study indicated higher hardness (2 to 3 times) compared to the hardness measured by the mixing law. In addition, the study demonstrated a very high compressive yield strength at both room temperature (RT) and > 1600°C[15]. This unusual behavior of RHEAs has drawn significant research attention to developing next-generation high-temperature materials with very high-temperature capability. Over the last decade, many RHEAs, based on Ti, Zr, Hf, V, Nb, Ta, Cr, Mo, and W refractory elements with/without adding non-refractory elements like Al, Co, Ni, and Si, have been studied with primary focus on high temperature (HT) strength, corrosion, and HT oxidation resistance[16–25]. The type and amount of alloying element can alter the characteristics of the material, described as the cocktail effect [26].

MoWTaNbV RHEAs, among the other RHEAs, have shown excellent potential for HT applications because of their promising properties and stability at HT for long durations as well as high specific strength[14, 27, 28]. This RHEA, however, demonstrates very limited RT ductility[15]. Studies indicate a compressive strain of ~ 1.7% (< 2%) before failure[15]. Han et al.[29] showed that the RT compressive ductility of MoWTaNbV increased with the addition of titanium (Ti) in equiatomic composition. The RT ductility increased from ~ 1.7% for MoWTaNbV to ~ 10.8 for MoWTaNbVTi. Another study[30] demonstrated the non-equiatomic composition of Ti also increases the RT ductility of MoWTaNb. The compressive ductility of MoWTaNb increased linearly with the increase of the Ti content. The increase in RT ductility due to Ti addition in RHEA could be attributed to the softening effect and the improvement of interstitial chemistry because of Ti addition. Although a minor drop in HT strength of MoWTaNbV due to Ti addition is observed (compressive yield strength of MoWTaNbV has decreased from 735 MPa to 659 MPa at 1200°C for MoWTaNbVTi), a significant improvement in RT ductility enables ease of manufacturing and processing of RHEA. Therefore, Ti could be a critical minor element in achieving the desired temperature performance and manufacturability of MoWTaNbV.

The RHEAs could achieve the performance target, but engineering difficulties of manufacturing these may bring minimum value to the structural applications. Historically, the primary limitation of refractory metal and alloys has been their manufacturability[5, 6]. The failure to achieve cost-effective manufacturing techniques for refractory metals and alloys widely limits their applications even though these may satisfy the performance matrixes[31]. The primary method of manufacturing, ingot metallurgy (melting/solidification followed by transformation of billet/plates, etc.) of refractory alloys, has been challenging and expensive mainly due to very high melting temperatures, low oxidation resistance, high hardness, and limited ductility[4, 32]. Initial studies of MoWTaNbV have been performed by manufacturing using the ingot metallurgy technique by vacuum arc melting (VAC)[15, 29, 33–35]. The dendritic structure, segregation, and relatively large microstructure, typical of casting characteristics, have been observed. The significant differences in melting points of the refractory elements aggravate the issues of microscopic defects like pores, macro segregations, and multiple phases in VAC MoWTaNbV[15, 29, 33–35]. Improving the microstructure characteristics may further require secondary mechanical processing, which will be difficult and expensive due to the relatively higher hardness and low ductility of RHEA. In addition, a machining step is required for MoWTaNbV applications. It is clear that the processing requirements of MoWTaNbV-based RHEA may outweigh the improved properties.

Powder metallurgy (PM) and related neat-net-shape (NNS) manufacturing methods have been potential cost-effective alternatives for manufacturing refractory alloys with controlled chemistry and microstructure characteristics. With the advancement of manufacturing methods for producing high-quality powders cost-effectively in bulk, the PM approach could be an ideal method for manufacturing RHEAs cost-effectively. In the PM approach, the consolidation of powders occurs before the melting of the refractory powders. Therefore, it requires relatively low energy for manufacturing compared to the ingot metallurgy methods. Many studies have reported the PM approach to manufacture MoWTaNbVTi RHEA[36–41]. The primary goal of the PM approach is to achieve a high density of consolidated powder close to the theoretical density. The majority of the PM approaches used the mechanical alloying of elemental powders to produce single-phase RHEA, followed by spark plasma sintering (SPS) to achieve the high density of RHEAs [36–41]. Long et al.[42], in their work, they have reported a significant improvement in compressive yield strength and compressive fracture strength at room temperature for NbMoTaWVTi RHEA when manufactured by mechanical alloying followed by spark plasma sintering compared to cast RHEA. The enhanced mechanical properties of PM NbMoTaWVTi could be attributed to the improved microstructure characteristics. The improvement in the microstructure and properties of MoWTaNbVTi compared to the VAC counterparts has been reported[37]. These studies indicate the potential of the PM approach for manufacturing RHEAs cost-effectively. Even though SPS can produce high-density RHEAs, it has limited applicability and is not economically favorable in the context of commercial viability and shape complexity.

The pressure-less sintering approach has the advantage of manufacturing near-net-shape products with relatively fewer manufacturing steps. The near-net-shape capability of manufacturing rather complex shapes in the sintering approach reduces the machining efforts. Due to process simplicity and cost considerations, the sintering approach has historically been the primary method for manufacturing refractory metals and alloys in laboratories and industries[43]. Over the years, a significant improvement in the sintering has been achieved to increase the density in many alloy systems, including refractory alloys. Using the additives and liquid phase sintering of the refractory metal and alloys has increased the sintering kinetics in these material systems, resulting in the increasing density of components and, therefore, the mechanical performance[44–46]. This has been attributed to the accelerated diffusion in refractory metals/alloys during sintering in the presence of minor additives/liquid phases[44, 47]. For instance, in W, `of transition elements in a minor quantity and enhanced sinterability compared to pure counterparts during sintering [48]. Near theoretical density of sintered W has been reported when minor amounts of Ni and Pd were added to the raw W powder[49]. A similar trend of increased density of the other refractory metals, Mo, Ta, and Nb, by additive transition elements has been observed[44]. Similar to refractory elements/alloys, the HEAs show sluggish diffusion characteristics[9, 50]. Sahil et al. [51]reported a high density of AlCoCrFeNi HEA using pressure-less sintering without additives and liquid phase considerations. However, the exact mechanism of densification of HEAs during sintering is not well understood.

The pressure-less sintering approach is successful for manufacturing refractory alloys[52–55], and an earlier study[56] on HEA indicates that powder-based sintering can manufacture RHEAs based on heavy refractory elements. The conventional pressure-less sintering of MoWTaNb RHEAs is rare in literature[57]. A recent study has shown the potential of manufacturing single-phase MoWTaNb RHEA using pressure-less sintering in SPS [58]. The evaluation of manufacturing efficiency for MoWTaNbVTi RHEA by the pressure-less sintering is interesting since this approach can potentially reduce the manufacturing cost of MoWTaNbVTi. The objective of this work is to demonstrate the capability of the pressure-less sintering method to manufacture heavy elements-based RHEAs in terms of density, phase formation, and elemental distribution. Various sintering parameters such as sintering temperature, time, composition, and addition of sintering enhancement agents have been investigated to explore the manufacturing efficiency for high density and understand the effect of sintering parameters on the sintered density, phase evolution, and elemental distribution of RHEAs. It is shown that the sintered density as high as > 99% can be achieved for MoWTaNbVTi_x RHEAs by pressure-less sintering. Also, the phase evolution during the manufacturing of RHEAs is of significant interest. Depending on the phase type in the microstructure, the characteristic properties of MoWTaNbVTi_x will be different[29, 37, 59]. Studies have demonstrated the different phase formation in MoWTaNbVTi_x depending on the manufacturing route[29, 35, 37]. Therefore, it is imperative to understand phase evolution in PM MoWTaNbVTi_x. In addition, elemental distribution in the microstructure of RHEAs has been a concern[36, 60–62]. Hence, this study also focuses on understanding the phase evolution and elemental distribution in MoWTaNbVTi_x manufactured via different PM approaches, pressureless sintering, HIP, and SPS.

RHEAs with varying compositions and processing routes were manufactured and characterized in the following ways. Elemental powders of Nb, Mo, Ta, W, V, Ti, and Ni with particle sizes in the range of -50 to 100 mesh size and purity of ≥ 99.9% were procured from Atlantic Equipment Engineers supplier for this study. These powders were mechanically alloyed (MA) in 500 ml tungsten vials using tungsten carbide balls with balls to powder ratio of 10:1 at 350 rpm in the presence of toluene. Three milling times of 15 hrs, 20 hrs, and 30 hrs were used. For 15 hrs and 20 hrs milled powders, refractory elements were added at the start of milling, and Ti/Ni were added 1 hr before the end of milling time, i.e., after 14 hrs for 15 hrs milled powders. For 30 hrs, milled powders Ti was added to the refractory elements at the start of milling (Ti blended). Small fractions of powders were taken out from the jar at the milling time intervals of 5 hrs, 10 hrs, 15 hrs, 20 hrs, and 30 hrs for x-ray diffraction (XRD).

Table 1

Manufactured MoWTaNbVTi_X compositions and the Alloy label.
Processing route	Alloy label	Ti (at. %)	Ni (wt% of MoWTaNbV)	Milling time	Sintering parameters	Sintering time
SSS	A1	8	0	15 hrs	1650ºC	8 hrs
LSS	A2	8	0	20 hrs	1700ºC	8 hrs
	A3	16	0	20 hrs	1700ºC	8 hrs
	A4	16 blended	0	30 hrs	1700 ºC	8hrs
AS	A5	8	1	20 hrs	1700ºC	8 hrs
LSS + HIP	A6	8	0	20 hrs	1800ºC + 200 MPa	4 hrs
LSS + HIP	A7	16	0	20 hrs	1800ºC + 200 MPa	4 hrs
SPS	A8	16	0	20 hrs	1700ºC + 50 MPa	5 mins
SPS	A9	27	0	20 hrs	1600ºC + 50 MPa	10 mins

MAed powders were transferred to a vacuum oven for drying at room temperature for 30 hrs in a vacuum of ~ 10^− 3 mbar. The dried powders were compacted into a cylindrical bar using a cold isostatic press (CIP) at a pressure of 310 Mpa for 5 mins. The green compacts were then sintered in a tube furnace using two different cycles. The sintering cycle schemes, solid state sintering (SSS) and liquid state sintering (LSS), are shown in Fig. 1a and Fig. 1b, respectively. The sintering temperatures for RHEAs were chosen based on the sintering temperature considered in the literature for the highest melting, W element. The sintering was performed at the maximum temperature of 1650°C and 1700°C following the two different sintering schemes for 8 hrs (a cumulative of 2 hrs at 1650°C and 6 hrs at 1600°C), and for 8hrs, respectively in argon (8 vol.% hydrogen) atmosphere. An initial dwell of 20 minutes at 300°C in the vacuum (10^− 4 bar) was added in both cycles to ensure completely dried powders were being sintered and also to remove the oxygen and moisture from the furnace. The final sintered samples were cylinders of ~ 7.6 cm in length and ~ 1.27 cm in diameter. List of all the samples manufactured are given in the Table 1. The selected sintered samples of A3 and A4 compositions were subjected to container less hot isostatic-pressing (HIP) to study the influence of HIP on sintered RHEAs. Activated sintering (AS) was done by addition of Ni to selected RHEA and sintered using LSS cycle. HIP was performed at Quintus Technologies AB, Västerås, Sweden, in a QIH21 HIP unit under a pressure of 200 MPa with a peak temperature of 1800ºC, and held at that pressure and temperature for 4hrs in Ar atmosphere (Fig. 1c). No significant change in the dimension of samples was observed after HIP processing of sintered RHEAs. Spark plasma sintering (SPS) was also used to manufacture bulk specimens with a diameter of 40 mm and thickness of 10 mm using 20 hr milled powders. Graphite dies were lined with graphite foil (GraFoil, NeoGraf Solutions, OH, USA) and coated with BN (ZYP Coatings Inc, TN, USA) to reduce adhesion and inhibit the carbiding of the alloying elements. Powders were loaded into the dies in an inert argon glovebox to avoid oxidation of the powders. Dies were loaded into an SPS (DCS-25, Thermal Technology, NV, USA) and evacuated to a vacuum of at least 2 x 10^− 2 Torr prior to heating. A load of 50 MPa was applied initially, and then specimens were heated via direct current resistive heating within the die. The heating profiles were 100°C/min to hold temperatures of either 1600°C (A11) or 1700°C (A12) and held at temperature for either 10 min (A11) or 5 min (A12), after which the load was removed and allowed to cool at 100°C/min to room temperature (Fig. 1d).

The cylindrical sintered samples were sectioned in disc shapes using electric discharge machining (EDM) for phase analysis and porosity measurement. X-ray diffraction (XRD) technique was done using a Rigaku SmartLab diffractometer with Cu Kα radiation at 40 kV and 40 mA and a 1-dimensional D/teX detector in conjunction with a Ni Kβ filter in 2θ range from 25° to 105°. Scans were performed at 6°/min with a 0.02°step size. The obtained XRD plots were analyzed using MDI Jade. Sintered RHEAs were polished and analyzed for SEM imaging using a Tescan Vega3 XMU scanning electron microscope equipped with an Ir SSD energy dispersive x-ray spectrometer. Mα radiation lines were used for elemental quantification and mapping of W and Ta since the accelerating voltage used for EDS analysis is 15 kV, which gives a lower overvoltage (< 2 for Ta, W) to characterize the minor quantities of Ta and W. The chemical analysis was performed using an electron probe micro-analyzer (EPMA) with a JEOL JXA8200 (Tokyo, Japan) electron microprobe equipped with five wavelength dispersive x-ray spectrometers (WDS) and an ultrathin-window energy dispersive spectrometer (EDS). Operating parameters were 12 kV accelerating voltage and 20 nA beam current. Analyses for light elements were accomplished with the WDS using multilayer synthetic crystals optimized for light elements Z = 4–7 (i.e. B through F). The LDE1 crystal (Layered Dispersive Element) has a d-space of 60 angstroms for measurement of O, and the LDE2 crystal has a d-spacing of 90 Å for measuring C. A Phi-Rho-Z program supplied with the third-party software Probe for EPMA (Probe Software, Eugene, OR) is used to reduce the raw data to elemental weight percent. The phase fractions were measured from the SEM micrographs with the aid of the image analysis software ImageJ. The micro hardness tester with a load of 1 kg and a hold time of 10 secs was used to measure the Vickers hardness. Ten hardness indents were considered from a sample to measure the average Vickers hardness values.

3.1. RHEAs Powder Characteristics:

The addition of Ti helps to improve the ductility and sinterability of RHEAs[63]. Ti and Ni powders were added 1hr before the end of milling to ensure that these elements do not alloy with matrix powders. The milling was continued after adding the Ti and Ni powders for another hour. The stepwise addition of Ni and Ti was done to prevent the formation of unwanted phases of Ni and Ti with refractory elements, which may adversely affect the properties of the RHEA[64]. Ti in the elemental form is assumed to increase sinterability by reacting locally with any interstitial contaminations[65]. Ni was added to the RHEA powder, assuming it would enhance the sinterability as it did in the refractory alloys[44]. The same procedure is employed to prepare the powder samples with different compositions except for alloy A4 (as shown in Table 1). For A4, Ti was added along with other refractory elements at the start of milling.

Figure 2a and Fig. 2b shows representative SEM micrographs of A1 alloy powders after 15 hr of mechanical milling. An irregular shape of particles is observed in the powders milled for 15 hr. This can be due to powder collisions during high-energy ball milling, leading to the plastic deformation of ductile powders and the fracture of brittle powders[66]. The brittle powders cold-weld with the ductile powders due to the ductile element’s plastic deformation, forming irregular morphology. The milling media also affects the powders' morphology [67]. For MoWTaNb powders, Tong et al. [67] reported the near-spherical shape of powders after milling when milled in acidic media (stearic acid) and lamellar morphology when milled in alkaline media (ethanol). In our study, irregular morphology of the powders milled in alkaline (toluene) media is consistent with the literature[67]. SEM micrograph at higher magnification (Fig. 2b) shows that few satellite particles are attached to the bigger particles, indicating the interaction of smaller particles with the other particles of RHEAs. The satellite particles are typical characteristics of mechanically milled powders of refractory metals and alloys when the particle sizes are in the submicron range. This is due to increase in the surface area of the particles with decrease in the particle size leading to agglomeration of the powder particles. Since the powders are lamellar/planar, the particle size is measured by determining the diameter of the imaginary spherical particle, equivalent to the particle's volume, using surface area, or projected area. The particle size is measured by measuring the diameter (${D}_{p}$) of the projected area (${A}_{pr}$) equivalent circle by using the following equation.

$${D}_{p}={\left(\frac{4{A}_{pr}}{\pi }\right)}^{\frac{1}{2}}$$

For a good approximation of overall particle distribution, a total of more than 300 particles are captured, as shown in Fig. 2a. The average particle size distribution is measured as 10.68 ± 9.46 µm, which is low compared to the average initial powder size (~ 80 µm), and a large standard deviation is observed. Although there is a relatively significant difference between the particle sizes, 90% of the particles were less than 20 µm, and almost 65% of the particles are less than 10 µm. Ti particles are ignored when calculating the average size and standard deviation as Ti was added 1hr prior to the end of milling time. The large particle size distribution is expected for the powders in a stage of mechanical activation as the equilibrium between the welding and fracturing of the powders is achieved after the powders reach the stage of mechanical alloying [67–69]. Hence, the particles are assumed to be in the stage of mechanical activation. The particle size distribution of the milled powder is shown in Fig. 2c. The particle size distribution is expected to vary with varying Ti content as Ti being the ductile powder among the brittle refractory powders will slow down the rate of decrease in particle sizes with milling time[70]. However, the time to mill for reaching the stage of mechanical activation will not vary significantly. Hence, the chosen 20 hr mill time will ensure that all RHEA powders are in the stage of mechanical activation. Many studies [36, 41, 60, 70, 71] indicated different characteristics of the powder shape and sizes. Liu et al. [46] reported that the particle size for 40 hr mill time is in the same range for MoTaNbVTi_x powders with more or less homogeneous particle sizes. Other studies have shown a similar trend of homogenized particle size after achieving a single phase of RHEAs during mechanical milling[36, 41, 60, 70, 71]. Peng et al. [72] for MoWTaNbV, Long et al. [37], and Liu et al. [73] reported the mill time of 40 hr for MoWTaNbVTi. Kang et al. [74] reported that 6 hr of mill time is required to completely alloy MoWTaNbVTi from the raw powder elements by milling in an argon atmosphere, while Pan et al. [75] reported a total mill time of 60 hr for the powders to be completely alloyed in acetone. Compared to the literature, the shorter milling time of 20 hr employed in this study reduces the contamination of powders from ball milling. The SEM images of the powders showed the difference in contrast between the powders, confirming that the powders were partially alloyed and are in the stage of mechanical activation.

Figure 3. shows the EDS analysis of 15 hr milled A1 powder. EDS was performed on selected particles based on the contrast of the particles. Carbon was eliminated from the EDS analysis as the powders were mounted on a carbon tap to avoid the error from the mounting. EDS measurement of 3 particles showed that some particles are rich in tungsten, some in molybdenum, and some in vanadium. Since a small fraction of local segregation in the powder microstructure is expected after 15 hr of the milling, and as EDS is performed on a small region instead of a whole particle, the vanadium-rich and molybdenum rich regions indicate the local vanadium and molybdenum segregations. The incomplete mechanical alloying observed in EDS also indicates the partial alloying of the powders. However, the homogenous equiatomic microstructure is expected to achieve after the sintering of the powders in the stage of mechanical activation[76].

Figure 4. illustrates the XRD patterns of A1 powders for total mill times of 20 min, 5 hr, 15 hr, and 20 hr and XRD patterns of A4 powders for total mill times of 20 hr and 30 hr. In the 20 min powder (initially blended powder), diffraction peaks corresponding to all pure elements are observed. Peaks corresponding to pure elements are observed to decrease with an increase in mill time, with peaks of low melting temperature elements decreasing faster than elements of high melting temperature, indicating slow diffusion of heavy refractory elements. A similar alloying sequence of low melting elements alloying first is reported for RHEAs in the literature[73]. After 20 hr of milling, the powders did not form a single phase, and the presence of peaks other than single BCC phase peaks in the XRD pattern indicates partial alloying. With the increase in mill time, the intensity of the peaks decreased, and a broadening of the peaks was observed because of an increase in lattice strain[77]. It was also observed that the peaks corresponding to V decreased faster and disappeared in 20 hr milled powders. However, for other elements with similar atomic radii ((Nb, Mo), (Ta, W)), it was hard to judge which element alloys faster based on the XRD pattern only. The effect of each element on the phase evolution was hard to estimate just by XRD pattern as RHEA was composed of multicomponent elements with similar atomic size and crystal structure. However, the novel idea of sequential alloying can be used to study the order of alloying[78].

After 20 hr of milling, the powders were not completely alloyed and were still in a stage of mechanical activation. Overlapping peaks were observed in 20 hr milled powders with higher peak broadening. These overlap peaks may correspond to metastable phases. During milling, the metastable phases can form[79]. The major peaks in the powders were at 40.28^o, 58.28^o, 73.28^o, and 87.14^o diffraction angles and correspond to a BCC crystal structure with a lattice constant of 3.165 Å. In 20 hr milled powders, peaks of V were observed to vanish, and peaks of Nb and Ta were significantly decreased, indicating that the lighter refractory element V tends to diffuse into the lattice of the heavier refractory elements (W and Mo). This phenomenon was expected as the XRD peaks of the primary phase in sintered samples were observed near the peaks of heavier refractory elements, which indicated the diffusion of lighter elements into the matrix of heavier elements during milling[75]. The high-intensity peaks in the RHEAs powders, which are assumed to be the primary phase, are observed to be near the peaks of W/Mo elements. Many studies have reported that the alloying sequence in the powders can be estimated based on the melting points of the elements[80–82]. W, being the highest melting point element in our RHEA composition, suggests that Ti, V, Nb, Mo, and Ta will dissolve into the W lattice[82]. Enthalpy of mixing can also be a contributing factor. The enthalpies of mixing for W with Nb, Mo, Ta, and V are negative. Negative enthalpy of mixing, therefore, can aid in the dissolution of Ti, V, Nb, and Ta atoms into the W lattice, as observed in the powder XRD patterns[81]. Hence, Ti, V, Nb, Mo, and Ta atoms dissolve into the W lattice during the milling and deforms the W lattice. The lattice constant of 3.165 Å for the BCC peaks in the powders is close to 3.165 Å, the lattice constant of W[83], further supporting that Ti, V, Nb, Mo, and Ta dissolve into the W matrix. However, we cannot conclude that the single BCC phase formed in the powders is the stable phase, as metastable phases can be formed in high energy ball milling [84].

The XRD pattern for 20 hr milled powder, when TiH₂ is added at the start of milling, and TiH₂ is added at the end, is observed to be the same. Hence, it can be assumed that Ti does not affect the alloying sequence in the RHEA powders. The absence of TiH₂ peaks in the powders is due to wt. % of TiH₂ being less than the XRD detection limit. But, in 20 hr milled powders, new overlapping peaks were observed near the peaks of (Nb, Ta). This might be because of an increase in the impurities from the milling media. These over lapping peaks are observed in both 20 hr milled powders irrespective of the sequence of Ti addition. In 30 hr milled blended powders, overlapping peaks were merged, and peaks of the secondary BCC phase were observed at 37.46º, 54.35º, 67.58º, and 80.06º diffraction angles with a lattice constant of 3.395 Å. The secondary BCC peaks were close to (Nb, Ta), indicating a secondary (Nb, Ta) rich BCC phase. The absence of Ti peaks in 30 hr milled blended powders suggests that Ti is dissolved into either a primary BCC phase rich in (Mo, W) or a secondary BCC phase rich in (Nb, Ta).

Multiple additional peaks were also observed in all powders but were not clear in 15 hr milled powders, indicating a small amount of the addition phase. These additional peaks were clearer in the powders of 20 hr, suggesting the evolution of a new phase with increased milling time. These peaks might correspond to the impurities introduced during high-energy energy milling. A similar study by Sarkar et al. [80] showed that peaks of WC are observed after 20 hr of milling RHEA powders in WC vials. The increase in the intensities of these peaks on 30 hr milled, blended powders suggests the increase in impurities in the RHEA powders with the increase in mill time. It is shown that impurities as high as 12% volume fraction were observed in single-phase ball-milled RHEA powders[80].

3.2. Sintered RHEAs Characteristics:

3.2.1. Phase evolution

XRD pattern of A1 (SSS) and A2, A3, A4 and A5 (LSS) are plotted and is shown in Fig. 5a. The diffraction peaks at 39.86º, 57.7º, 72.44º, 86.12º corresponding to (110), (200), (211), and (220) planes, respectively, of a BCC structure (BCC_LSS) with a lattice parameter of 3.193 Å has been observed in A1. However, the major BCC phase peaks in LSS RHEAs are observed to be shifted towards the right with diffraction peaks at 40.00º, 58.02º, 72.98º, and 86.80º corresponding to (110), (200), (211), and (220) planes, respectively, of a BCC structure (BCC_LSS) with a lattice parameter of 3.176 Å close to reported SPS processed RHEAs[37]. The shift in BCC peaks of LSS RHEAs might be due to an increase in solubility of smaller radius Ti atoms in the BCC matrix in the liquid state sintered RHEAs as the diffusion is higher in LSS[47]. However, the position of BCC peaks is not affected by the amount of Ti in the LSS RHEAs, indicating that the solubility limit of Ti in the matrix is reached in the A2 with 8 at.% of Ti. Hence, it is safe to assume that excess Ti precipitates out in high Ti containing A3 and A4. It is also observed that peak intensities of the matrix BCC phase in LSS RHEAs decreased with an increase in Ti content, further supporting the argument of Ti precipitation.

Interestingly, a greater fraction of the BCC refractory matrix phase in SSS compared to LSS RHEAs, as evident by the higher peak intensities of the matrix BCC phase in SSS A1 than in LSS RHEAs is observed. This suggests that during LSS, a decrease in the volume fraction of the BCC matrix phase occurred, which might be due to some of the refractory components precipitating into a different phase. The variation in the BCC phase composition in both RHEAs may cause variation in the lattice parameter of the BCC phase in LSS and SSS RHEAs. This may be due to the difference in diffusion characteristics of elements in both LSS and SSS processes. The lattice parameter of the matrix BCC phase in the sintered samples is observed to increase from 3.165 Å in the powders to 3.176 Å in the LSS sintered samples. Since the powders were not completely alloyed, the bigger radius elements of (Nb, Ta) were not completely dissolved into the matrix. When sintered, the solubility of (Nb, Ta) in the BCC matrix increases due to enhanced diffusion in LSS. Hence, there is an increase in the lattice parameter of the primary BCC phase.

In addition to the BCC peaks, minor FCC peaks (FCC_1) were observed in all RHEAs. The FCC peaks are more clearly visible in A3 and A4 with peaks at 35.38º, 41.10º, 59.56º, 71.28º, and 74.98º with a lattice parameter of 4.385 Å. These FCC peaks were matching closely with niobium carbide (NbC – PDF # 04-002-0729) and titanium niobium tantalum carbide ((TiTaNb)_0.51C_0.49 – pdf # 04-012-9728). The higher intensity FCC peaks indicate a higher FCC phase fraction in A3 and A4 than in A1 and A2. Multiple oxide peaks were also observed in all RHEAs except in A4. In A4, peaks corresponding to TiO_0.637 (PDF # 04-005-4341) were observed. Since the Ti was blended initially, it was uniformly distributed in the RHEA powder, which possibly helped the formation of one type of Ti-O phase during sintering. This indicates that more than one Ti-O phase was possibly formed in all RHEA samples except in A4.

Additionally, some minor peaks corresponding to carbide were also observed. These carbides were probably formed due to the contamination of powders from the WC milling media. The addition of Ni in A5 for activated sintering showed no change in the XRD pattern compared to A2. This implies that Ni has a limited impact on the phase evolution during the sintering of RHEAs. Hence, it is safe to assume that sintering is completed in the LSS RHEAs, and the addition of Ni has a limited effect on the microstructure evolution of RHEAs.

Compared to powders, the sintered XRD pattern of RHEAs showed sharp and high-intensity peaks. X-ray diffraction theory says that the position of the diffraction peak and peak width depends on the lattice constant and grain size of the grains[85]. During high-energy ball milling, the powder mixture undergoes large strain, causing peak broadening. However, elevated temperatures promote diffusion during sintering, alleviating the stresses and causing grain growth. As a result, in contrast to the diffraction patterns of alloy powder, the diffraction peaks of alloys become sharper, and the diffraction intensity increases [36]. Figure 5b compares the XRD pattern of HIPed A6 and A7 with LSS A2 and A3 respectively. Minor variations in peak intensities can be observed after the HIPing. Similar to LSS RHEAs, a major BCC_LSS peaks, a minor FCC peaks with some oxide and carbide peaks can be observed.

Figure 5c compares the XRD pattern of SPS A8 and A9 with A3. Like LSS RHEAs, A8 showed major BCC_LSS and FCC_1 along with oxide and carbide peaks. Unknown peaks observed in the range of 48º to 50º in LSS RHEAs disappeared in SPS-processed RHEAs. However, oxide and carbide peaks are observed in A8, similar to A3, suggesting a similar microstructure in both A3 and A8. In A9, FCC peaks (FCC_2) were observed to move right with peaks at 36.06º, 41.92º, 60.78º, and 76.58º with a corresponding lattice constant of 4.307 Å. FCC_2 peaks matched closely with khamrabaevite (TiC – PDF # 03-065-8805). The shift in FCC peaks might be due to the difference in the composition of the carbide phase. Since A9 RHEA was sintered at 1700ºC, a liquid phases SPS process is expected. Due to high Ti content (~ 27 at.%), high diffusion rates, and shorter sintering time, it is possible for a fraction of free liquid Ti to react with carbon and oxygen impurities in the powder and form TiC and TiO_x avoiding the formation of complex carbide and oxide phases.

The micrographs of sintered RHEAs were captured using SEM and shown in Fig. 6. Near full density is achieved in all processes despite the sluggish diffusion in RHEAs. A three-phase microstructure is observed in all RHEAs with dark and grey precipitated phases in a white matrix. However, the grey phase is noticeably higher in high Ti containing A3, A4, A7, A8 and A9. It is safe to assume that the primary white matrix phase is a BCC phase based on XRD. The three-phase fractions were measured using ImageJ [86] and reported in Table 2. The volume fraction of the dark phase is observed to increase with an increase in Ti content. This suggests that the dark phase might be rich in Ti. No pores were observed in the matrix of all RHEAs, suggesting that complete sintering of RHEAs is achieved in pressureless sintering and SPS processes. However, shrinkage pores in the dark phase were observed in A1, A2 and A3 (Fig. 6a, Fig. 6b and Fig. 6c respectively) when Ti was added later in milling. The observation of pores in the dark phase suggests that the dark phase is a low melting phase. During the sintering of multi-phase powders, phases with melting temperatures less than the peak sintering temperature melt and promote the sintering. Since Ti was added 1 hr before the end of milling, it is safe to assume that Ti is not dissolved in other phases at the end. As the sintering temperature in the LSS cycle is higher than the melting temperature of Ti, Ti particles in milled powders melt, enhancing the diffusion and improving the sinterability. However, larger liquid-forming particles can create pores when solidified during cooling[87]. Since Ti was added 1hr before the end of milling, the Ti particles were bigger than the RHEA powders in all the samples. The liquid-forming Ti particles were comparatively bigger. Liquid Ti particles volume increased during sintering and shrank during solidification, forming pores[88]. No pores were observed in A4 (Fig. 6d). As Ti was added at the start of milling and Ti particles were comparatively smaller in milled powders, no shrinkage pores were observed in A4 [47].

In the LSS process, the microstructure is dominated by the solubility behavior. In a microstructure with high densification, the solubility of the liquid-forming phase in the solid is usually low, and the solubility of the solid phase in the liquid-forming phase is high[87]. Because of the high solidification observed in RHEAs, it can be assumed that the liquid-forming Ti particles were not soluble in the solid BCC matrix, and hence, liquid-forming Ti particles are precipitated as the Ti-rich dark phase.

Figure 6e and Fig. 6f show the SEM micrographs of A6 and A7, respectively, after HIP processing. The three-phase microstructure is retained after HIPing, showing the microstructural stability of the sintered RHEAs. However, HIP processing of sintered RHEAs helped eliminate shrinkage pores in the dark phase. It is known that pressure-assisted sintering will enhance diffusion and help with the homogenization of microstructure and elimination of pores [89]. The phase fractions of dark and grey phases in pressureless sintering have remained the same after HIP processing. Since Ti is precipitated as the dark phase in sintered RHEAs, there is a possibility of forming a liquid phase during the HIP processing of RHEAs as the peak sintering temperature in the HIP process is 1800ºC, above the melting temperature of Ti. At the peak sintering temperature, the Ti-rich liquid phase will experience additional pressure and fill the spaces between the particles. This filling under pressure helps eliminate shrinkage pores, and relative density can reach ~ 100% [90].

Figure 6g and Fig. 6h show the SEM micrographs of SPS processed A8 and A9 RHEAs, respectively. Similar to pressureless sintered and HIP-processed RHEAs, a three-phase microstructure is also observed in A8 and A9. The absence of pores indicates a near full-density solid formation after SPS. Li et al. [91] showed that porosity could be eliminated in SPS-processed MoTaVW when sintered to a temperature of 1800ºC. The addition of Ti in the present RHEAs helped decrease the peak sintering temperature to 1650ºC to eliminate pores. This can be attributed to the relatively high Ti diffusivity that helps increase the sinterability of the RHEAs powders[92].

Previous studies have shown precipitated phases in SPS-processed RHEAs even though XRD of the powders indicated a single phase[36–40, 61, 62, 71–75, 91, 93–95]. However, the precipitated phase in the same SPS-processed RHEA is inconsistent in the literature. Most previous studies showed precipitation of Nb/Ta as Ta/Nb/(Ta, Nb) rich precipitated phase in SPS processed RHEAs [36, 71, 72, 74, 75, 93]. Long et al.[71] showed precipitation of Ta₂VO₆ tetragonal phase in SPS processed MoWTaNbVCr RHEA. You et al.[93] reported precipitation of (Ti,Nb) rich phases in SPS processed W_yTaNbVTi_x RHEAs. Peng et. al.[72] indicated the precipitation of (Ta,V)O₂ particles in SPS-processed MoWTaNbV. Pan et al.[75] reported (Nb, Ta) rich precipitated phase. Kang. et al. reported precipitation of (Ta,V) rich precipitated phase in SPS processed MoWTaNb RHEA.

The difference in the sintered RHEAs can be due to the difference in the impurities introduced during manufacturing processes. It is well known that impurities are unavoidable in mechanical alloying, and the amount and type of impurities depend on the milling conditions. Since refractory elements have a high affinity towards interstitials like C and O, common impurities in ball milling, they tend to react with them and form precipitated phases. A detailed literature study (Table 3) revealed that the precipitated phases in PM-processed RHEAs depend on the milling conditions. For example, studies that used WC balls and vials for MA of powders, similar to the present study, reported (Ta, Nb) rich carbide phase after SPS processing[39, 40, 75, 94]. Since the milling media is the primary source of impurities in the MA process, WC is the dominant impurity when WC balls and vials are used. The WC impurities can either precipitate or break down and act as the carbon source in the presence of elements with a high affinity towards carbon.

Table 2

Phase fractions of three phases in RHEAs
RHEA	Bright	Grey	Dark
A1	86.80 ± 1.65	5.86 ± 0.90	7.89 ±0.89
A2	84.46 ± 1.84	7.77 ± 1.02	7.83 ±1.29
A3	70.51 ± 1.50	14.01 ± 1.13	15.58 ± 1.47
A4	68.48 ± 0.35	15.49 ± 0.50	16.18 ± 0.32
A5	77.88 ± 0.38	14.32 ± 0.38	7.79 ± 2.24
A6	84.10 ± 1.82	7.24 ± 1.41	8.65 ± 0.99
A7	69.39 ± 0.77	16.62 ± 0.53	14.02 ± 0.83
A8	74.02 ± 2.57	14.06 ± 2.67	11.98 ± 0.69
A9	69.99 ± 2.78	14.43 ± 1.81	15.62 ± 1.84

Table 3

Precipitated phases in RHEAs were reported in the literature.
RHEA	Milling media	Balls material	Vial material	Precipitated phase in sintered RHEA
MoNbTaTiV[36]	Ar	Stainless steel	Stainless steel	FCC-TiO/TiN
NbMoTaWVCr[71]	Ar	Stainless steel	Stainless steel	Tetragonal Ta₂VO₆
NbMoTaWV[96]	Glove box	Stainless steel	Stainless steel	Tetragonal (Ta,V)O₂
NbMoTaWVTi[97]	Ar	Stainless steel	Stainless steel	FCC – (Ti,O) rich
MoNbTaTiV[38]	Ar	Stainless steel	Stainless steel	FCC –phase
MoNbTaTiW[61]	Ar + stearic acid	Hardened steel	Hardened steel	FCC – (Ti) rich
WNbMoTaV[60]	Ar	Stainless steel	WC	Tetragonal - Ta₂VO₆
WTaMoNbV[62]	Ar	WC	Stainless steel	Tetragonal - Ta₂VO₆
NbMoTaW[98]	stearic acid	Stainless steel	-	FCC – (Nb,Ta) rich
MoNbTaW[39]	Ar + stearic acid	WC	WC	FCC – (Nb,Ta)C rich
WNbMoVTa[40]	stearic acid	WC	WC	HCP – (Ta, Nb)C rich
NbMoTaWRe[94]	Ar	WC	WC	FCC – (Nb,Ta)C rich
Ti₈NbMoTaW[41]	Ar + acetone (PCA)	WC	WC	BCC – (Ta,Nb) rich

3.2.3 Elemental distribution

Representative elemental maps of A3 and A7 RHEAs are obtained by EPMA analysis and are shown in Fig. 7a and 7b, respectively. Refractory elements were concentrated mainly in the white matrix region with minute concentrations of (Ti,O). The precipitated dark phase is observed to be rich in Ti and O. Among all the elements in RHEA powders, Ti has a high affinity for O [99]. Hence, Ti might have preferentially picked up the O during manufacturing and precipitated as the oxide phase during the manufacturing cycle, i.e., MA and sintering. This helps reduce O content in the matrix phase, which is highly desirable. Also, it is observed that the grey phase has higher concentrations of (Nb, Ta, Ti, C) with minor concentrations of (V, O) and almost zero concentrations of group VI elements (Mo, W). This indicates that the grey phase is a (NbTaTi) C-rich phase. The observation of the grey phase in sintered RHEAs might be due to the preferential reaction of (NbTaTi) with C, as the milling media is rich in C. The higher Ti content in the grey phase may be due to the higher solubility of Ti in Nb and Ta than in Mo and W[64]. It is to be noted that the XRD pattern of the sintered RHEAs showed only two phases with carbide and oxide peaks, while a three-phase microstructure is observed in all RHEAs.

Table 4 indicates the average chemical composition of the three phases, matrix, dark, and grey, in A3 and A10 RHEAs. The similarity in the chemical composition of A3 and A7 supports the argument that the sintered microstructure is stable. The group VI refractory elements (Mo,W) are higher than the initial mixture composition in the bright matrix phase, while (V,Nb,Ta) were less than the initial mixture observed in both RHEAs. In both cases, the Ti was < 2 at.% in the matrix phase and precipitated out from the matrix phase into two different phases, dark (~ 40 at.%) and grey (~ 15 at.%) phases. The dark phase is Ti-O rich, while the grey phase is (NbTaTi)C rich. A large standard deviation in oxygen and Ti at. % are observed in the dark phase of both A3 and A7. It is possible to form multiple TiO_x dark phases with varying x in sintered RHEAs as Ti was not dissolved into the matrix at the start of sintering. Hence, a large standard deviation for Ti and O is observed in the dark phase in these samples. This also suggests that multiple oxide peaks observed in the XRD pattern of the RHEAs can be from different TiO_x peaks. This explains the difference in the observation of two phases in the XRD pattern, while a three-phase microstructure is observed in SEM micrographs of sintered RHEAs. The XRD peaks corresponding to oxide peaks may be due to the dark phase TiO_x. Based on the EPMA chemical analysis, the composition of the grey phase was determined to be (TiTaNb)_0.53C_0.47 with < 1 at.% of Mo and W. Hence the observed FCC peaks correspond to (TiTaNb)_0.51C_0.49, carbide in the XRD pattern of the sintered RHEAs. The C impurities entered during manufacturing react with Nb, Ta, and Ti to form a carbide of FCC crystal structure (FCC_1) with a lattice parameter of 4.385 Å.

Some theoretical studies have shown that MoNbTaWTi forms a single-phase BCC alloy at high temperatures [15, 29, 33–35]. However, precipitation of (Ti,O) rich and (Nb,Ta,Ti,C) rich phases were observed in this study. Hence, impurities like C and O can significantly affect the microstructure of RHEAs and affect the formation of precipitates (Nb,Ta,Ti,C) rich and (Ti,O) rich phases. It is interesting to note that FCC peaks that are observed in sintered RHEAs were absent in powders. This may suggests absence of these precipitate phases in powders after milling.Stress released during the sintering of powders could increase elemental diffusion and phase precipitation. Hence, introducing C and O during milling can promote preferential precipitation of carbide and oxide phases in sintered RHEAs. Since Ti has a higher affinity towards O than C, it reacts preferentially with O and forms a dark precipitated phase rich in (Ti,O) with minor quantities of C. With a higher affinity of Nb > Ta > Ti with carbon, a grey phase rich in (Nb,Ta,Ti,C) with minor quantities of V was formed. Carbon mainly comes from WC milling balls and vials and a small fraction from the decomposition of toulene[100]. Kannunikova et al. [100] reported TiC_xH_y carbohydride formation in the mechanical activation process of Ti in toluene. Liu et al.[39] has shown that WC can decompose to form (Ta,Nb) carbide as Gibbs free energy of formation of (Ta,Nb) carbide is lower than (W,Mo) carbide[101]. Liu et al. [39] reported a theoretical study using Scheil Gulliver cooling simulations and showed that Ta, Nb, and Ti have a higher tendency to combine with C.

Table 4

Elemental compositions of RHEAs A3 and A7.
RHEA	Phase	C	O	Ti	V	Nb	Mo	Ta	W
A3	Bright	7.33 ± 1.29	2.63 ± 0.84	1.56 ± 0.08	16.60 ± 0.42	16.37 ± 0.72	17.55 ± 1.01	15.87 ± 0.46	22.08 ± 1.287
	Grey	44.35 ± 1.14	2.65 ± 0.42	14.12 ± 1.47	2.67 ± 0.56	16.98 ± 0.73	0.59 ± 0.62	18.01 ± 0.51	0.63 ± 0.78
	Dark	11.70 ± 1.04	45.52 ± 3.39	41.17 ± 2.29	0.56 ± 0.08	0.57 ± 0.05	0.06 ± 0.01	0.34 ± 0.03	0.08 ± 0.01
A7	Bright	5.11 ± 0.18	1.51 ± 0.42	0.95 ± 0.33	17.92 ± 0.94	16.19 ± 0.43	20.07 ± 0.41	15.18 ± 0.45	23.06 ± 0.40
	Grey	42.94 ± 1.53	4.29 ± 2.47	14.66 ± 3.15	3.59 ± 0.20	17.06 ± 2.31	0.33 ± 0.12	16.87 ± 2.00	0.28 ± 0.14
	Dark	4.68 ± 0.71	53.76 ± 3.07	40.47 ± 2.34	0.93 ± 0.32	0.10 ± 0.05	0.01 ± 0.01	0.03 ± 0.01	0.02 ± 0.01

3.2.5 Effect of Ni on phase distribution

Many researchers have shown that Ti and Ni can be used as sintering aids to enhance the sinterability of refractory elements[44, 102–108]. It has proven to be effective for processing high-density heavy refractory elements and alloys as it can improve the overall diffusion, hence decreasing the sintering temperature and time [109]. Such improvement in atomic diffusion is important in RHEAs as RHEAs tend to exhibit sluggish diffusion, which can lower the sinterability. Shivakumar et al.[109] showed that adding Ni to MoWTaNb can improve the densification behavior in sintered RHEAs. This is attributed to the activated sintering by forming the liquid phase in MoWTaNb RHEA. Since Ti[37] and Ni[109] have very low solubility in RHEAs and were added 1 hr before the end of milling, Ni in liquid form is expected to increase the overall liquid phase and enhance the density of A5 during sintering. According to the binary phase diagram of Ti-Ni, the eutectic point for Ti-Ni is 942ºC at 28.4 wt.% Ni, lowering its melting temperature [110]. During the heating phase in the sintering cycle, when the temperature exceeds the melting points of Ni and Ti, the liquid phases of Ni and Ti are formed. The formation of the liquid phase with solid particles around it leads to an interface between liquid and solid phases. As the sintering proceeds, the amount of liquid phase increases gradually, and viscosity decreases, which can increase the flowability of liquid between the particles. The particles will come closer to each other through the force developed due to the difference in liquid surface tension of the dual phases. This force will help decrease the additional interfacial energy developed by reducing the surface area through the coarsening liquid phase. Therefore, the addition of Ni lowers the liquid forming temperature during sintering and provides additional time for particle rearrangement.

Figure 8. are the SEM micrographs of the sintered A2 and Ni containing A5. Similar to A2, a three-phase microstructure is observed with similar amounts of carbide phase in Ni containing A5 (Table 2). This indicates that Ni is not participating in the precipitation of carbide phase. However, it is observed that spherical Ti-O phase fraction and diameter (as observed as a circle in 2D view of the micrographs in Fig. 8b) increased. Phase fractions of Ti-O precipitate increased to 16.77 ± 0.33 in Ni containing A5 from 8.69 ± 0.38 in Ni-free A2 with the same manufacturing cycle and Ti content. Also, the diameter of the spherical Ti-O phase increased to 5.31 ± 0.32 in A5 from 3.37 ± 0.21 in A2. These observations indicate that Ni promotes Ti diffusion in RHEAs during sintering by forming low-temperature eutectic liquid with Ti before RHEAs sintering starts. The increased volume of the liquid phase in Ni containing A5 ensures a larger volume of liquid solidifies after the sintering cycle, leading to a larger fraction and diameter of the Ti-rich Ti-O phase.

3.2.6. Hardness

Figure 9 shows the hardness of RHEAs with varying Ti content and processing routes. The hardness varied in the range of ~ 690–911 Hv. HIP-processed A6 showed the highest hardness, and activated sintered A5 showed the lowest. Among A3, A4, A7, and A8 RHEAs with 16 at.% of Ti, SPS-processed A8 showed the highest hardness, and LSS A3 showed lower hardness. A slight increase in the hardness of A4 was observed when Ti was added at the start of milling compared to A3. Among the four different processing routes, activated sintered A5 exhibit lower hardness, and SPS-processed RHEAs exhibit higher hardness. The observed hardness is much higher than the reported values of 786 Hv for MoWTaNbVTi_x in the literature[29, 97]. The higher hardness of PM MoWTaNbVTi_x in the present study can be attributed to the carbide phase in the microstructure.

The variation in hardness of the MoWTaNbVTi_x RHEAs is due to varying amounts of precipitated (Ti,O) rich dark oxide phases with low hardness (~ 317 Hv)[111] and (TiTaNb)_0.53C_0.47 carbide phase with high hardness of [112] compared to hardness of matrix phase[14]. A comparison of the hardness of RHEAs with respective oxide and carbide phase fractions is summarized in Table 5. Since the porosity of the RHEAs is very low, the effect of porosity on hardness is expected to be low, and we can safely ignore it The hardness behavior of the RHEAs can be understood by considering both oxide and carbide phase fractions simultaneously. Among the RHEAs A1, A2, A5, and A6 with 8 at.% of Ti, A5 has the highest oxide-to-carbide phase fraction ratio, followed by A1. This ratio is smaller A6 explaining the higher hardness. The same argument can be applied to RHEAs A3, A4, A7, and A8 with 16 Ti at.%. The order of A3 ~ A4 < A7 < A8 oxide carbide ratio follows the measured hardness trend of A3 ~ A4 < A7 < A8. High hardness in SPS samples might be due to lower TiO_x soft phases. Figure 9 shows the nonlinear variation in hardness as a function of oxide-to-carbide ratio. Hardness behavior indicates that it primarily depends on oxide to carbide ratio and is not influenced by the at. % of Ti and PM processing method. This indicates that contamination is the leading factor in controlling the hardness of the PM processed RHEAs.

Table.5 Comparison of Vickers hardness with oxide and carbide phase fractions in RHEAs.

RHEA	Hardness (Hv)	Oxide phase fraction (%)	Carbide phase fraction (%)
A1	689.9 ± 46.8	7.89 ±0.89	5.86 ± 0.90
A2	861.5 ± 48.6	7.83 ±1.29	7.77 ± 1.02
A3	755.5 ± 42.4	15.58 ± 1.47	14.01 ± 1.13
A4	809.4 ± 39.4	16.18 ± 0.32	15.49 ± 0.50
A5	602.1 ± 35.5	16.77 ± 0.33	7.24 ± 1.41
A6	909.0 ± 55.9	8.65 ± 0.99	7.29 ± 1.08
A7	842.8 ± 48.7	14.02 ± 0.83	16.62 ± 0.53
A8	911.2 ± 35.6	11.98 ± 0.69	14.06 ± 2.67
A9	835.9 ± 28.5	15.62 ± 1.84	14.43 ± 1.81

NbMoTaWVTi_x refractory high entropy alloy was manufactured by various PM processes, including pressureless sintering (solid state, liquid state, and activated), hot isostatic pressing (solid state and liquid state), and spark plasma sintering (solid state and liquid state). It is shown that all PM processing routes, including the conventional pressureless sintering method, can achieve high-density sintered RHEAs. Titanium oxide phases with varying oxygen concentrations can form in Ti added sintered MoWTaNbV RHEAs in the PM manufacturing approach, largely due to the affinity of Ti with oxygen. It was found that Ti precipitates into two phases, as (TiTaNb)_0.53C_0.47 and TiO_x, in the non-equiatomic MoWTaNbV matrix of PM manufactured MoWTaNbVTi_x despite the PM processing routes. The fraction of the carbide and TiO_x phases in the microstructure depends on the amount of Ti and PM processing method of MoWTaNbVTi_x. A significant variation is observed in the hardness of PM MoWTaNbVTi_x. The variation in the hardness among different PM processed and Ti composition MoWTaNbVTi_x is attributed to the variation in the ratio of oxide to carbide phase fractions in the microstructure. A higher oxide-to-carbide phase ratio shows a lower hardness and vice-versa. Adding Ni as an activator helps enhance Ti diffusion during the sintering MoWTaNbVTi_x, while a limited effect of diffusion of heavy elements was observed. The mechanical activation approach can be employed to reduce the milling time, thereby decreasing the impurities in the milled powders and sintered bulk MoWTaNbVTi_x RHEAs.

Present work concludes that pressureless sintering can be used to produce high-density RHEAs. However, contamination of the powders during milling can be a severe problem, leading to the formation of precipitated phases. The differences in the precipitate phases observed in PM manufactured RHEAs across the literature are addressed and attributed to the differences in the impurities in the milled powders. Sintered RHEAs showed no change in phase composition after HIP processed, showing the microstructure stability and potential candidate as a high-temperature creep-resistant material. A non-linear relation of hardness with oxide to carbide ratio is observed, indicating that impurities play a major role in the hardness behavior of PM MoWTaNbVTi_x RHEAs.

Competing Interests

The authors have no relevant financial or non-financial interests to disclose

Funding

This work was supported by NASA EPSCoR through New Mexico Space Grant Consortium (Grant Number 80NSSC21M017).

Author Contributions

Material preparation was performed by Surya Bijjala, Ryan Wilkerson and Chad Beamer. Data collection and analysis was performed by Surya Bijjala and Pankaj Kumar. The first draft of the manuscript was written by Surya Bijjala and all authors commented on previous versions of the manuscript. All authors read and approved the final manuscript. Funding acquisition was done by Pankaj Kumar.

Acknowledgment

STB and PK greatly acknowledge the financial support from the NASA EPSCoR through New Mexico Space Grant Consortium (Grant Number 80NSSC21M017).

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Understanding the phase evolution and elemental distribution in MoWTaNbVTix manufactured via powder metallurgical approach

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Abstract

Figures

1. Introduction

2. Experimental procedures

3. Results and discussions

3.1. RHEAs Powder Characteristics:

3.2. Sintered RHEAs Characteristics:

3.2.1. Phase evolution

3.2.3 Elemental distribution

3.2.5 Effect of Ni on phase distribution

3.2.6. Hardness

4. Conclusion

Declarations

Competing Interests

Funding

Author Contributions

Acknowledgment

References

Supplementary Files

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