High Temperature Performance of Spark Plasma Sintered W 0.5 (TaTiVCr) 0.5 Alloy

: The phase stability, compressive strength, and tribology of tungsten alloy containing low activation elements, W 0.5 (TaTiVCr) 0.5 , at elevated temperature up to 1400 ◦ C were investigated. The spark plasma sintered W 0.5 (TaTiVCr) 0.5 alloy showed body centered cubic (BCC) structure, which was stable up to 1400 ◦ C using in-situ high temperature XRD analysis and did not show formation of secondary phases. The W 0.5 (TaTiVCr) 0.5 alloy showed exceptionally high compressive yield strength of 1136 ± 40 MPa, 830 ± 60 MPa and 425 ± 15 MPa at 1000 ◦ C, 1200 ◦ C and 1400 ◦ C, respectively. The high temperature tribology at 400 ◦ C showed an average coe ﬃ cient of friction (COF) and low wear rate of 0.55 and 1.37 × 10 − 5 mm 3 / Nm, respectively. The superior compressive strength and wear resistance properties were attributed to the solid solution strengthening of the alloy. The low activation composition, high phase stability, superior high temperature strength, and good wear resistance at 400 ◦ C of W 0.5 (TaTiVCr) 0.5 suggest its potential utilization in extreme applications such as plasma facing materials, rocket nozzles and industrial tooling.


Introduction
Tungsten (W) and its alloys possess high melting point, good mechanical properties at elevated temperature, high hardness, low activation in radiation environment and low sputtering yield, which makes it a potential material for various high temperature and nuclear applications [1]. Furthermore, W exhibits low coefficient of thermal expansion, good thermal conductivity and low vapor pressure [2][3][4]. The advantages of W are coupled with its shortcomings, such as its low fracture toughness, radiation-induced embrittlement, blistering at moderate temperatures by Deuterium (D) and Helium (He), and formation of pits, holes and bubbles by Helium at higher temperatures [5][6][7]. In order to overcome the shortcomings in W, new strategies on adding different alloying elements to W, such as W-Re [8,9], W-Ta [10,11], W-V [12,13], W-Mo [14,15], W-Cr [16][17][18] and W-Ti [19,20] binary alloys, W-based composites [21][22][23][24][25], or nanostructure-engineered W [26][27][28] have been investigated. The reported W-based binary alloys have shown significant improvement in the mechanical properties [19,20]. However, experimental analysis in different aspects of binary alloys have revealed numerous constraints as well, such as irradiation induced embrittlement in W-Re and W-Os, deterioration of mechanical properties due to metastable phases in W-Ti and W-V alloys [1,9]. Similarly, nanocrystalline W has been found to generate He bubble under irradiation leading to poor mechanical properties at high temperatures [27,29]. The laminated, fiber-and net-reinforced W-based

Structural and Chemical Characterization
X-ray diffraction (XRD) analysis was carried out using Cu-Kα radiation in PANalytical Empyrean operating at 40 kV and 40 mA. Scans were performed between 2θ range of 5 and 100 degrees with 10 mm divergence slit and 1 • diffracted beam slits. For high temperature in-situ XRD analysis, HT furnace was installed with W-heating strip. The sample was placed onto a W-heating strip and a heating rate of 10 • C/min was used to increase the temperature up to 1400 • C, where measurements were performed at every 100 • C interval after a holding time of 30 min and measurement time of 40 min. The high temperature in-situ XRD analysis was carried out in vacuum atmosphere (10 −4 mbar) using Anton Paar HTK 16N strip heater (PANalytical Empyrean, Malvern, UK). The density of the sintered sample was measured using Archimedes' method. In order to carry out selected-area electron diffraction (SAED) pattern analysis, a thin lamella (~50 nm) was prepared by Focused Ion Beam (FIB) milling using Helios Nanolab 600 Dual Beam System (Field Electron and Ion Company, Hillsboro, Oregon, OR, USA), and examined under Transmission Electron Microscope (JEM2100F, JEOL Ltd., Tokyo, Japan). Surface and wear track morphology of the sintered alloy was carried out using scanning electron microscopy (SEM, JSM-IT300LV, JEOL GmbH, Freising, Germany) and energy dispersive X-ray spectroscopy (EDS) with an accelerating voltage of 10-20 kV and working distance of 10 mm.

High Temperature Mechanical and Tribological Characterization
High temperature compression tests were carried out at 1000 • C, 1200 • C and 1400 • C using Gleeble-3800 (DSI, New York, NY, USA) in vacuum environment with a heating rate and strain rate of 10 • C/s and 10 −3 s −1 , respectively. For high temperature compression tests, cylindrical bars of 4 mm diameter and 6 mm in length were cut using wire electrical discharge machining (EDM) to obtain samples with an aspect ratio of 1.5. Two tests per temperature were performed for repeatability. Tribological tests were performed with a ball-on-disc setup at 400 • C in air atmosphere using universal tribometer (Rtec, San Jose, CA, USA). W 0.5 (TaTiVCr) 0.5 sample with 12 mm diameter and 3 mm height was used against a counterball of Al 2 O 3 with a diameter of 9.5 mm. The wear tests were carried out at 5 N normal load with a sliding speed of 0.1 m/s for 1 h corresponding to a sliding distance of 600 m. Three tests were performed to examine the repeatability of wear rate and coefficient of friction (COF). The wear depth was calculated using optical profilometry (Wyko NT1100, Bruker Inc., Billerica, MA, US). The wear rate was calculated by the ratio of wear volume to sliding distance and normal load.

Results and Discussion
The phase formation and microstructure of W 0.5 (TaTiVCr) 0.5 alloy sintered at 1600 • C is shown in Figure 1. The X-ray diffractogram of the sintered alloy showed formation of BCC solid solution with a lattice parameter of 0.314 nm, and formation of TiC phase (Figure 1a). TiC phase forms due to the reaction of liquid Ti with the graphite die during the spark plasma sintering [42]. The relative density of the sintered sample was measured to be 14.1 g/cm 3 . The microstructure of the sintered alloy showed formation of three different phases consisting of HEA (grey region), W-rich (bright region) and TiC (black region) phase ( Figure 1b). The EDS elemental maps of the alloying elements in the sintered alloy show distribution of elements in W-rich, HEA-phase (marked in red circles in corresponding micrograph), and Ti-rich phase (TiC), in Figure 1c. The formation of HEA-phase (grey region) was further confirmed using EDS point analysis from grey region and bright region, marked as 1 and 2 in Figure 1b respectively, as presented in Table 1.   The chemical nature of various regions (W-rich, TiC and HEA) in W 0.5 (TaTiVCr) 0.5 alloy was investigated using transmission electron microscopy (TEM). The TEM-EDS and selected-area electron diffraction (SAED) patterns of different regions are shown in Figure 2. The TEM-EDS elemental maps showed uniform distribution of all elements except Ti-rich zone and HEA phase, as shown in Figure 2a  The in-situ high temperature XRD analysis of W0.5(TaTiVCr)0.5 alloy from RT to 1400 °C is shown in Figure 3. The BCC solid solution phase was observed to be stable at all temperatures without the formation of intermetallic phases ( Figure 3a). The stability of the alloy up to 1400°C can be related to high melting point of W, where the composition consists of more than 50 at. % W, which aids thermal stability. Furthermore, the high concentration of other elements enhances the thermal stability through lattice distortion leading to low thermal diffusion in the sintered alloy and enhancing the stability of BCC phase. The formation of TiO2 was observed after in-situ XRD at 1200 °C, which can be related to softening of Ti leading to its diffusion to the surface and presence of small amount of oxygen in the environment that reacts with Ti on the surface, as shown in Figure 3c and its respective EDS analysis (spot 2) in Table 2 [48]. The SEM-EDS after in-situ high temperature XRD analysis showed the W-rich phase and HEA phase to be stable after the measurement, however the oxide formation in all the phases is observed, as shown in Figure 3c and Table 2. The exact influence of each element at high temperature towards diffusion and oxide formation has yet to be determined and requires further investigation [49].  The in-situ high temperature XRD analysis of W 0.5 (TaTiVCr) 0.5 alloy from RT to 1400 • C is shown in Figure 3. The BCC solid solution phase was observed to be stable at all temperatures without the formation of intermetallic phases (Figure 3a). The stability of the alloy up to 1400 • C can be related to high melting point of W, where the composition consists of more than 50 at. % W, which aids thermal stability. Furthermore, the high concentration of other elements enhances the thermal stability through lattice distortion leading to low thermal diffusion in the sintered alloy and enhancing the stability of BCC phase. The formation of TiO 2 was observed after in-situ XRD at 1200 • C, which can be related to softening of Ti leading to its diffusion to the surface and presence of small amount of oxygen in the environment that reacts with Ti on the surface, as shown in Figure 3c and its respective EDS analysis (spot 2) in Table 2 [48]. The SEM-EDS after in-situ high temperature XRD analysis showed the W-rich phase and HEA phase to be stable after the measurement, however the oxide formation in all the phases is observed, as shown in Figure 3c and Table 2. The exact influence of each element at high temperature towards diffusion and oxide formation has yet to be determined and requires further investigation [49]. The in-situ high temperature XRD analysis of W0.5(TaTiVCr)0.5 alloy from RT to 1400 °C is shown in Figure 3. The BCC solid solution phase was observed to be stable at all temperatures without the formation of intermetallic phases (Figure 3a). The stability of the alloy up to 1400 °C can be related to high melting point of W, where the composition consists of more than 50 at. % W, which aids thermal stability. Furthermore, the high concentration of other elements enhances the thermal stability through lattice distortion leading to low thermal diffusion in the sintered alloy and enhancing the stability of BCC phase. The formation of TiO2 was observed after in-situ XRD at 1200 °C, which can be related to softening of Ti leading to its diffusion to the surface and presence of small amount of oxygen in the environment that reacts with Ti on the surface, as shown in Figure 3c and its respective EDS analysis (spot 2) in Table 2 [48]. The SEM-EDS after in-situ high temperature XRD analysis showed the W-rich phase and HEA phase to be stable after the measurement, however the oxide formation in all the phases is observed, as shown in Figure 3c and Table 2. The exact influence of each element at high temperature towards diffusion and oxide formation has yet to be determined and requires further investigation [49].   The engineering stress-strain curves from high temperature compression tests are shown in Figure 4. The sintered alloy showed the highest yield strength of 1136 ± 40 MPa at 1000 • C, followed by yield strength of 830 ± 60 MPa at 1200 • C, and finally yield strength of 425 ± 15 MPa at 1400 • C (Figure 4b). The high temperature compressive strength of W 0.5 (TaTiVCr) 0.5 alloy was found to be more than twice that of pure W (~200 MPa at 1350 • C) [50], and higher than previously reported RHEA of NbMoTaW (548 MPa, 506 MPa and 421 MPa at 1000 • C, 1200 • C and 1400 • C, respectively) [51], TaNbHfZrTi (295 MPa and 92 MPa at 1000 • C and 1200 • C, respectively) [52], NbTaTiV (437 MPa at 1000 • C) [53], TaNbVTiAl (~360 MPa and 142 MPa at 1000 • C and 1200 • C, respectively) [54], and VCrFeTaW (182-371 MPa at 1000 • C) [55] due to solid solution strengthening and precipitation strengthening (from in-situ formed TiC phase) [42,56]. The decrease in the yield strength with increasing temperature can be related to the diffusion of Ti, as shown in Figure 3c [49]. The ductility was observed to increase from 6% to 10% with increasing temperature due to thermal softening [57], and Ti diffusion due to its low melting point. Furthermore, the yield strength of W 0.5 (TaTiVCr) 0.5 alloy at all temperatures was found to be 3-4 times higher than the previous reported W-based heavy alloys [58,59].
Metals 2020, 10, x FOR PEER REVIEW 6 of 12 Table 2. EDS analysis of W0.5(TaTiVCr)0.5 alloy in Figure 3 after in-situ XRD measurement at 1400 °C in at. %. The engineering stress-strain curves from high temperature compression tests are shown in Figure 4. The sintered alloy showed the highest yield strength of 1136 ± 40 MPa at 1000 °C, followed by yield strength of 830 ± 60 MPa at 1200 °C, and finally yield strength of 425 ± 15 MPa at 1400 °C (Figure 4b). The high temperature compressive strength of W0.5(TaTiVCr)0.5 alloy was found to be more than twice that of pure W (~200 MPa at 1350 °C) [50], and higher than previously reported RHEA of NbMoTaW (548 MPa, 506 MPa and 421 MPa at 1000 °C, 1200 °C and 1400 °C, respectively) [51], TaNbHfZrTi (295 MPa and 92 MPa at 1000 °C and 1200 °C, respectively) [52], NbTaTiV (437 MPa at 1000 °C) [53], TaNbVTiAl (~360 MPa and 142 MPa at 1000 °C and 1200 °C, respectively) [54], and VCrFeTaW (182-371 MPa at 1000 °C) [55] due to solid solution strengthening and precipitation strengthening (from in-situ formed TiC phase) [42,56]. The decrease in the yield strength with increasing temperature can be related to the diffusion of Ti, as shown in Figure 3c [49]. The ductility was observed to increase from 6% to 10% with increasing temperature due to thermal softening [57], and Ti diffusion due to its low melting point. Furthermore, the yield strength of W0.5(TaTiVCr)0.5 alloy at all temperatures was found to be 3-4 times higher than the previous reported W-based heavy alloys [58,59]. The fracture surface morphologies of specimens after high temperature compression tests are shown in Figure 5 and EDS point analysis from different locations, spot 1-7, marked on fracture surfaces in Figure 5 are presented in Table 3. The fracture surface of the specimen fractured at 1000 °C during compression showed inter-granular fracture around W-rich phase, as shown in Figure 5 a,b. The high strength at 1000 °C can be related to the presence of solid solution strengthening and TiC phases (marked as red arrows in Figure 5b and EDS spot 2 in Table 3), which can significantly enhance the impediment of crack propagation [60,61]. A total strain of 0.1 was observed at 1000 °C compression test, higher than total compression strain at RT observed in earlier work [42]. Similar fracture mode was observed at 1200 °C along with a slightly more softening, as shown in Figure 5 c,d. However, Ti-rich zone was observed using EDS at grain boundary, as shown in EDS spot 4 in Table 3. A further increase in the ductility and wetting behavior were observed for the alloy at 1400 °C (Figure 5e,f), which can be due to further softening of HEA-phase, as shown in EDS spot 6 and 7 in Table 3 [62]. The fracture surface morphologies of specimens after high temperature compression tests are shown in Figure 5 and EDS point analysis from different locations, spot 1-7, marked on fracture surfaces in Figure 5 are presented in Table 3. The fracture surface of the specimen fractured at 1000 • C during compression showed inter-granular fracture around W-rich phase, as shown in Figure 5a,b. The high strength at 1000 • C can be related to the presence of solid solution strengthening and TiC phases (marked as red arrows in Figure 5b and EDS spot 2 in Table 3), which can significantly enhance the impediment of crack propagation [60,61]. A total strain of 0.1 was observed at 1000 • C compression test, higher than total compression strain at RT observed in earlier work [42]. Similar fracture mode was observed at 1200 • C along with a slightly more softening, as shown in Figure 5c,d. However, Ti-rich zone was observed using EDS at grain boundary, as shown in EDS spot 4 in Table 3. A further increase in the ductility and wetting behavior were observed for the alloy at 1400 • C (Figure 5e,f), which can be due to further softening of HEA-phase, as shown in EDS spot 6 and 7 in Table 3 [62].  Table 3. EDS analysis on fracture surface in Figure 5 in at. %. The high temperature sliding wear test of W0.5(TaTiVCr)0.5 alloy at 400 °C against alumina counterball was performed to study the wear behavior before the start of oxidation, as shown in Figure 6. The sliding wear test showed an average COF of 0.55 after an initial run-in period up to 200 m of sliding distance (Figure 6a). The alumina counterball showed a wear scar diameter of 2 mm arising from high wear induced by W0.5(TaTiVCr)0.5 alloy, as shown in Figure 6b. The EDS analysis on alumina counterball showed mostly W transfer layer (spot 1 in Table 4). The wear track morphology of W0.5(TaTiVCr)0.5 alloy is shown in Figure 6c,d. The sliding wear showed mild adhesive  Table 3. EDS analysis on fracture surface in Figure 5 in at. %. The high temperature sliding wear test of W 0.5 (TaTiVCr) 0.5 alloy at 400 • C against alumina counterball was performed to study the wear behavior before the start of oxidation, as shown in Figure 6. The sliding wear test showed an average COF of 0.55 after an initial run-in period up to 200 m of sliding distance (Figure 6a). The alumina counterball showed a wear scar diameter of 2 mm arising from high wear induced by W 0.5 (TaTiVCr) 0.5 alloy, as shown in Figure 6b. The EDS analysis on alumina counterball showed mostly W transfer layer (spot 1 in Table 4). The wear track morphology of W 0.5 (TaTiVCr) 0.5 alloy is shown in Figure 6c,d. The sliding wear showed mild adhesive wear behavior and high amount of alumina counterball transfer, as shown in EDS data of spot 2 and 3 in Table 4. The sliding wear of W 0.5 (TaTiVCr) 0.5 alloy against alumina counterball resulted in an average wear rate of 1.37 × 10 −5 mm 3 /Nm, which was found to be three times lower than the wear rate of pure tungsten (5 × 10 −5 mm 3 /Nm) at 500 • C as reported by Jiang et al. [63]. Recently, Zhou et al. developed a (FeCoCrNi) 1−x (WC) x (x = 3-11, at. %) composite HEA using SPS to increase the wear properties of FCC based HEA [64]. The study showed enhancement of wear resistance up to 5 at. % WC addition against GCr15 steel counter ball. However, the wear resistance decreased for higher WC content due to lower bonding between reinforcement and HEA matrix phases. In contrast, formation of in-situ TiC reinforcement in HEA matrix has been found to be better for the wear resistance [65][66][67]. The formation of in-situ carbide reinforcement enhances the mechanical and tribological properties due to stronger bonding between the ceramic reinforcement and metal matrix [68]. Similarly, in this work, formation of in-situ TiC reinforcement, as marked in Figure 6d, coupled with the high hardness (788 HV [36]) from BCC solid solution in W 0.5 (TaTiVCr) 0.5 alloy enhanced the wear resistance even at 400 • C.

Spot
Metals 2020, 10, x FOR PEER REVIEW 8 of 12 wear behavior and high amount of alumina counterball transfer, as shown in EDS data of spot 2 and 3 in Table 4. The sliding wear of W0.5(TaTiVCr)0.5 alloy against alumina counterball resulted in an average wear rate of 1.37 × 10 −5 mm 3 /Nm, which was found to be three times lower than the wear rate of pure tungsten (5 × 10 −5 mm 3 /Nm) at 500 °C as reported by Jiang et al. [63]. Recently, Zhou et al. developed a (FeCoCrNi)1−x(WC)x (x = 3-11, at. %) composite HEA using SPS to increase the wear properties of FCC based HEA [64]. The study showed enhancement of wear resistance up to 5 at. % WC addition against GCr15 steel counter ball. However, the wear resistance decreased for higher WC content due to lower bonding between reinforcement and HEA matrix phases. In contrast, formation of in-situ TiC reinforcement in HEA matrix has been found to be better for the wear resistance [65][66][67]. The formation of in-situ carbide reinforcement enhances the mechanical and tribological properties due to stronger bonding between the ceramic reinforcement and metal matrix [68]. Similarly, in this work, formation of in-situ TiC reinforcement, as marked in Figure 6d, coupled with the high hardness (788 HV [36]) from BCC solid solution in W0.5(TaTiVCr)0.5 alloy enhanced the wear resistance even at 400 °C.  The high temperature studies of W0.5(TaTiVCr)0.5 alloy showed high temperature stability of BCC phase up to 1400 °C, high yield strength up to 1400 °C compared to previously reported values, and superior wear resistance at 400 °C. The resulting behavior of W0.5(TaTiVCr)0.5 alloy at high temperature shows its promising application in aerospace and fusion plasma facing material. In order to ensure safety in applications of W0.5(TaTiVCr)0.5 alloy, it is required to study its resistance to irradiation, oxidation and erosion as well, which are the topics of our forthcoming studies.  Table 4. EDS analysis on counterball and wear track in Figure 6 in at. %. The high temperature studies of W 0.5 (TaTiVCr) 0.5 alloy showed high temperature stability of BCC phase up to 1400 • C, high yield strength up to 1400 • C compared to previously reported values, and superior wear resistance at 400 • C. The resulting behavior of W 0.5 (TaTiVCr) 0.5 alloy at high temperature shows its promising application in aerospace and fusion plasma facing material. In order to ensure safety in applications of W 0.5 (TaTiVCr) 0.5 alloy, it is required to study its resistance to irradiation, oxidation and erosion as well, which are the topics of our forthcoming studies.

Conclusions
In this work, we have studied the elevated temperature performance of spark plasma sintered W 0.5 (TaTiVCr) 0.5 alloy. The study of high temperature phase stability using in-situ XRD revealed stable BCC solid solution up to 1400 • C without formation of intermetallic phases. However, formation of TiO 2 was observed after 1200 • C due to higher diffusion rate of Ti compared to other elements at high temperatures. The high temperature compression tests showed high yield strength of 1136 ± 40 MPa at 1000 • C, 830 ± 60 MPa at 1200 • C and 425 ± 15 MPa at 1400 • C due to the presence of in-situ formed TiC phase and solid solution strengthening in W-rich phase and HEA phase. High temperature tribological test at 400 • C showed superior wear resistance due to combination of in-situ TiC phase and HEA solid solution. The wear loss in the alloy was caused mostly by high amount of transfer layer from alumina counterball.

Conflicts of Interest:
The authors declare no conflict of interest.