Refractory High-Entropy Hftatinbzr-Based Alloys by Combined Use of Ball Milling and Spark Plasma Sintering: Eﬀect of Milling Intensity

$Refractory High-Entropy Hftatinbzr-Based Alloys by Combined Use of Ball Milling and Spark Plasma Sintering: Eﬀect of Milling Intensity$

metals

Article Refractory High-Entropy HfTaTiNbZr-Based Alloys by Combined Use of Ball Milling and Spark Plasma Sintering: Eﬀect of Milling Intensity

Natalia Shkodich 1,2,*, Alexey Sedegov 1, Kirill Kuskov 1 , Sergey Busurin 3, Yury Scheck 2, Sergey Vadchenko 2 and Dmitry Moskovskikh 1

1 Center of Functional Nanoceramics, National University of Science and Technology MISIS, Moscow 119049, Russia; [email protected] (A.S.); [email protected] (K.K.); [email protected] (D.M.) 2 Merzhanov Institute of Structural Macrokinetics and Materials Science, Russian Academy of Sciences, Chernogolovka, Moscow 142432, Russia; [email protected] (Y.S.); [email protected] (S.V.) 3 NTO IRE-POLUS, LCC, Fryazino, Moscow 141190, Russia; [email protected] * Correspondence: [email protected]; Tel.: +7-49652-46-299

Received: 1 September 2020; Accepted: 18 September 2020; Published: 20 September 2020

Abstract: For the ﬁrst time, a powder of refractory body-centered cubic (bcc) HfTaTiNbZr-based high-entropy alloy (RHEA) was prepared by short-term (90 min) high-energy ball milling (HEBM) followed by spark plasma sintering (SPS) at 1300 ◦C for 10 min and the resultant bulk material was characterized by XRD and SEM/EDX. The material showed ultra-high Vickers hardness (10.7 GPa) and a density of 9.87 0.18 g/cm3 (98.7%). Our alloy was found to consist of HfZrTiTaNb-based ± solid solution with bcc structure as a main phase, a hexagonal closest packed (hcp) Hf/Zr-based solid solution, and Me2Fe phases (Me = Hf, Zr) as minor admixtures. Principal elements of the HEA phase were uniformly distributed over the bulk of HfTaTiNbZr-based alloy. Similar alloys synthesized without milling or in the case of low-energy ball milling (LEBM, 10 h) consisted of a bcc HEA and a Hf/Zr-rich hcp solid solution; in this case, the Vickers hardness of such alloys was found to have a value of 6.4 GPa and 5.8 GPa, respectively.

Keywords: refractory high-entropy alloys; high-energy ball milling; low-energy ball milling; spark plasma sintering; bcc solid solution

1. Introduction The classical approach to designing alloys, based on the use of one or two principal elements and minor amounts of other elements added for property enhancement, has been extended by Yeh et al. [1] and Cantor [2], who independently suggested a new concept for the design of multicomponent alloys, termed high-entropy alloys (HEAs). Newly developed HEAs containing at least five components in equiatomic or nearly equiatomic amounts (ranging between 5 and 35 at. %) have received much attention in view of their unique structures and excellent properties [3–9]. The experimental data currently available for more than 400 different HEAs unanimously confirm their unique mechanical, chemical, and electrical properties [10]. The presence of several principal elements in HEAs not only characterizes their composition, but also brings about fundamental physical changes in their configurational entropy, free energy, phase selection, and stability [10,11]. In general, the high-entropy effect in these alloys can greatly reduce the free energy of simple solid-solution phases, especially at high temperatures, thus making chemically ordered intermetallic compounds less competitive owing to the formation of simple solid solutions with a face-centered cubic (fcc), body-centered cubic (bcc), hexagonal closest packed (hcp), or their combinatorial structures [1,2,10].

Metals 2020, 10, 1268; doi:10.3390/met10091268 www.mdpi.com/journal/metals Metals 2020, 10, 1268 2 of 15

The formation of solid solutions in HEAs is governed by some thermodynamic parameters such as entropy of mixing (∆Smix), enthalpy of mixing (∆Hmix), and atomic size difference (δ)[10]. According to [12], HEAs form a simple solid solution when 22 ∆H 7 kJ/mol, 11 ∆S 19.5 kJ/mol, − ≤ mix ≤ ≤ mix ≤ and δ 8.5%. ≤ Moreover, it has been found that a crucial factor, which defines whether or not fcc or bcc solid solutions appear are formed in HEAs, is the valence electron concentration (VEC) [13]. It has been observed that bcc and fcc phases are formed for VEC < 6.87 and VEC 8.0, respectively, while both of ≥ the phases (i.e., bcc + fcc mixture) coexist for 6.87 VEC 8.0 [13]. ≤ ≤ The first single-phase CoCrFeMnNi material (also known as the Cantor alloy) has an outstanding combination of properties, such as exceptional resistance to damage and resistance to destruction, especially at cryogenic temperatures [14–16]; high tensile strength and ductility [3–5,17]; superplasticity [5]; and high hardness and strength [3,6,7]. Typical synthesis routes used for the fabrication of traditional alloys can also be applied to the synthesis of HEAs. On the basis of a type of intermixing procedure, existing processing routes can be classified into four categories [18–21]: from liquid phase (arc melting), from solid state (mechanical alloying), from gas phase (sputtering), and from electrodeposition process. It has been found that different synthesis routes applied to the same starting compositions produce materials with distinctly different microstructure and properties. Thus, the CoCrFeNiTi alloy synthesized by arc melting presents a fcc–bcc mixture with minor admixture of intermetallic phase [15], while that prepared by mechanical alloying (MA) has a single fcc phase [17]. Over the last few years, ever growing attention has been given to HEAs based on refractory metals such as Cr, Hf, Mo, Nb, Ta, Ti, V, W, and Zr [10,22–24]. Meanwhile, these materials having high potential for high-temperature applications still remain to be studied adequately. Refractory TaNbHfZrTi HEA (bcc solid solution, Hv = 3826 MPa, σ0.2 = 929 MPa) was produced [23] by arc melting and subsequent consolidation by hot pressing at T = 1473 K and p = 207 MPa for 3 h. MoNbTaTiV refractory high-entropy alloy with ultra-fine and homogeneneous microstructure (bcc solid solution) was successfully fabricated by mechanical alloying for 40 h followed by spark plasma sintering (SPS) at 1500–1700 ◦C[25]. Mechanical alloying (MA) is a well-known solid-state, non-equilibrium, top-down approach for the fabrication of nanocrystalline materials. It involves milling elemental powders to achieve mixing on an atomic scale. This technique has been successfully applied to produce a variety of materials, like intermetallics, ordered compounds, solid-solution alloys, amorphous structures, quasi-crystalline phases, and nanocomposites [26–29]. In addition to nanoscale processing, MA also takes advantage of extended solubility of solids even in immiscible systems [26,30]. This is attributed to elevated diffusion rates arising as a result of a nanosize of starting powder blends. Besides increased configurational entropy, MA imparts enhanced stability to solid-solution phases in HEAs. The first synthesis of HEAs by MA is credited to Varalakshmi et al. [31], who prepared nanocrystalline AlCrCuFeTiZn alloy back in 2008. Since then, MA has been progressively used to synthesize nanocrystalline HEAs, as evidenced by a growing number of publications on mechanically alloyed HEAs. Most solid-solution phases in HEAs form within 15–40 h of milling. Prolonged milling hours are often used for the production of amorphous HEAs [27]. For consolidation of HEAs powders, conventional sintering methods are often used, which include vacuum hot pressing sintering, microwave sintering, and hot isostatic pressing. Among these, a highly promising technique for sintering and production of HEA materials is spark plasma sintering (SPS) [32–34]. The main difference between SPS and conventional hot pressing is that, in SPS, the current is applied as many pulses over very short times (less than milliseconds (ms)) instead of one single pulse for a long time. As with most methods where pressure and temperature are applied simultaneously, both pressure and temperature are lower than would be required if applying them in subsequent steps. This means that lower sintering temperatures and shorter holding times are required, and as a result, less grain growth and better properties can be obtained [33,35]. Metals 2020, 10, 1268 3 of 15

An equiatomic WNbMoTaV high entropy alloy (HEA) with a bcc structure was ﬁrst fabricated by the powder metallurgical process of mechanical alloying (MA) and spark plasma sintering (SPS). MA process, microstructure, and mechanical properties of the WNbMoTaV alloy were studied systematically. During MA, a single bcc phase was formed and its mean particle size and crystallite size were estimated as 1.83 µm and 66.1 nm, respectively, after 6 h of MA. In this communication, we report on the preparation of equiatomic HfTaTiNbZr-based refractory high-entropy alloy (RHEA) with a bcc structure by combined use of short-term high-energy ball milling and spark plasma sintering and on their structural and mechanical characterization.

2. Experimental HfTaTiNbZr powder mixtures were prepared by two different pathways: (a) low-energy ball milling (LEBM) and (b) high-energy ball milling (HEBM) of the following elemental powder blends: Hf (99.1% pure, 100–200 µm), Ta (99.9%, 40–60 µm), Ti (99.2%, 40–60 µm), Nb (99.9% pure, 40–60 µm), and Zr (99.6%, 100–200 µm)—taken in equiatomic amounts. Mechanical treatment was performed in a water-cooled double-station planetary ball mill Activator-2s (Activator, Russia) using stainless-steel cylindrical jars and balls (5–7 mm in diameter). In all cases, the ball/powder ratio was 20:1 (by weight). The vial was evacuated and then filled with Ar gas at 4 atm. Ball milling was run at different rotation speeds of sun wheel/jars: 200/400 rpm (LEBM) and 694/1388 rpm (HEBM). Milling time (t) was varied between 20 and 90 min in the case of HEBM and between 1 and 10 h in the case of LEBM. Starting and milled HfTaTiNbZr powder mixtures were SPS-consolidated in a vacuum in a Labox 650 facility (Sinter Land, Niigata, Japan). Powder mixture was placed into a cylindrical graphite die (inner diameter 15 mm) and uniaxially compressed at 50 MPa. Because one of HEA phases can be expected to exist at temperatures above 1020 ◦C[36], the sample was heated at a rate of 100 ◦C/min up to some preset sintering temperature in the range 1100–1700 ◦C by passing rectangular current pulses through the sample. The dwell time (τ) at sintering temperature was 10 min. SPS-produced disks were 2–3 mm thick and 15 mm in diameter. Structure/composition of initial, milled, and consolidated powders were characterized by XRD (DRON-3M diffractometer, Saint Petersburg, Russia, Cu-Kα radiation, 2θ = 20–95◦), SEM (Vega 3 TESCAN, Brno, Czech Republic), and EDX (Oxford Inca spectrometer, Oxford, UK) using Aztec software. Microhardness was measured using an Emco-Test DuraScan 70 apparatus (Kuchl, Austria) under a load of 9.8 N. The thermal stability of initial and RHEA powder mixtures was studied by DSC (STA 449 F1 Simultaneous Thermal Analyzer, Netzsch, Germany) up to 1500 ◦C with a heating rate of 20 ◦C/min in an Ar atmosphere.

3. Results and Discussion For predicting the formation of solid solutions in the HfTaTiNbZr system, we applied the well-established concepts of the HEA community [10] based on the Hume–Rothery rules, that is, taking into account composition-weighted terms for diﬀerences in atomic radii (∆r), electronegativity (∆χ), and average valence electron concentration (VEC) and using the data presented in Table1.

Table 1. Properties of the elements under consideration [8] (see the text for details). hcp, hexagonal closest packed; bcc, body-centered cubic; fcc, face-centered cubic; VEC, valence electron concentration; EN, electronegativity.

Elements Atomic Number Structure Radius, pm Tm (K) VEC Pauling EN Hf 72 hcp 157.75 2233 4 1.30 Ta 73 bcc 143 3017 5 1.50 Ti 22 hcp 146.15 1670 4 1.54 Nb 41 bcc 142.9 2468 5 1.60 Zr 40 hcp 160.25 1851 4 1.33 Metals 2020, 10, 1268x FOR PEER REVIEW 4 of 16 15

Nb 41 bcc 142.9 2468 5 1.60 It has been observedZr that40 the bcc hcpand fcc phases160.25 in the1851 system 4 under 1.33 study are formed for VEC < 6.87 and VEC 8.0, respectively, while both phases (i.e., bcc + fcc mixture) coexist when 6.87 ≥ ≤ VECFor8.0 the [13 HfTaTiNbZr]. composition, the atomic size difference (Δr = 6.06%) is well within the ≤ adoptedFor limits the HfTaTiNbZr of 0 ≤ Δr ≤ 8.5%. composition, VEC = 4.4 the falls atomic into the size range difference for the (formation∆r = 6.06%) of a is bcc well solid within solution the (VECadopted < 6.87) limits [13]. of 0The∆ electronegativityr 8.5%. VEC = 4.4(EN) falls difference into the is range 0.1445. for the formation of a bcc solid solution ≤ ≤ (VECAs< 6.87)known [13 [12],]. The electronegativity electronegativity Δχ (EN) has dilittlefference or no is influence 0.1445. on the formation of solid-solution or amorphousAs known phase. [12], electronegativity Even when no ∆rangeχ has has little been or no prescribed influence for on thethe formationformation of of solid-solution the disordered or solidamorphous solution, phase. it was Even noted when [3 no7] rangethat a has larger been value prescribed of Δχ for is theexpected formation to aid of thethedisordered formation solidof a compound.solution, it was In our noted case, [37 Δχ] that = 0.1176. a larger value of ∆χ is expected to aid the formation of a compound. In our case,In the∆χ present= 0.1176. study, the following three approaches were applied: (a) SPS of elemental powdersIn the (Hf, present Ta, Ti, study, Nb, Zr), the following(b) combination three approaches of LEBM and were SPS, applied: and (c) (a) combination SPS of elemental of HEBM powders and (Hf,SPS. Ta, Ti, Nb, Zr), (b) combination of LEBM and SPS, and (c) combination of HEBM and SPS.

3.1. SPS SPS of Elemental Powders Initial powder blends were prepared from elemen elementaltal powders of of Hf, Hf, Ta, Ta, Ti, Nb, and Zr taken in equatomic amountsamounts andand consolidated consolidated by by SPS SPS at at di ffdifferenterent temperatures temperatures (1100 (1100◦C, ° 1300С, 1300◦C, 1500°С, 1500◦C, and°С, 1700and 1700◦C) in°С order) in order to produce to produce HfTaTiNbZr HfTaTiNbZr RHEA. RHEA. Before sintering, the blended po powderswders exhibited diffraction diffraction line liness from all of the constituent elements. SPS SPS at at 1100 °C◦C is is seen seen (Figure (Figure 11)) toto resultresult inin peaks’peaks’ broadening,broadening, theirtheir overlapping,overlapping, andand thethe appearance ofof aa signal signal at at 33.56 33.56°◦ corresponding corresponding to theto the Zr/Hf-based Zr/Hf-based solid solid solution. solution. The di Theffraction diffraction peaks peaksfrom initial from elementsinitial elements are still are present still present in the XRDin the pattern. XRD pattern.

Figure 1. XRD patterns of initial HfTaTiNbZr powder blend and HfTaTiNbZr bulk specimens sintered Figure 1. XRD patterns of initial HfTaTiNbZr powder blend and HfTaTiNbZr bulk specimens at different temperatures (indicated). HCP, hexagonal closest packed; BCC, body-centered cubic. sintered at different temperatures (indicated). HCP, hexagonal closest packed; BCC, body-centered cubic.On the basis of the SEM/EDX results (Figures1 and2), we may conclude that the formation of different solid solutions is due to material melting at grain boundaries. SPS at 1300 ◦C was found to yieldOnbcc thesolid basis solution of the asSEM/EDX a main phase.results However,(Figures 1 theand presence 2), we may of broadconclude XRD that peaks the indicatesformation the of differentincomplete solid dissolution solutions ofis metals.due to material The diff ractionmelting pattern at grain in boundaries. Figure2 also SPS shows at 1300 the °C presence was found of Hf- to yieldand Zr-based bcc solid solid solution solutions. as a main phase. However, the presence of broad XRD peaks indicates the incomplete dissolution of metals. The diffraction pattern in Figure 2 also shows the presence of Hf- and Zr-based solid solutions. The XRD results well agree with phase composition and EDX maps (Figure 2) for the HfTaTiNbZr sample consolidated at 1300 °C (dwell time 0 min). The principal elements are shown in color-scale contours. Pure Zr and Hf are still present in the bulk sample.

Metals 2020, 10, 1268 5 of 15

Metals 2020, 10, x FOR PEER REVIEW 5 of 16

FigureFigure 2. Representative 2. Representative compositional compositional EDX mappingEDX mapping results results for the for HfTaTiNbZr the HfTaTiNbZr sample consolidatedsample consolidated at 1300 °C (τ = 0 min). at 1300 ◦C(τ = 0 min).

The XRDUpon results SPS at 1500 well °C, agree the withcrystallinity phase compositionof the bcc phase and grew, EDX although maps (Figure the diffraction2) for the peaks HfTaTiNbZr from Zr andFigure two 2.solid Representative solutions based compositional on Zr and HfEDX (hcp mapping with a = results3.245, cfor =5.140 the andHfTaTiNbZr bcc with asample = 3.455) are sample consolidated at 1300 ◦C (dwell time 0 min). The principal elements are shown in color-scale still consolidatedpresent in the at 1300form °C of ( τlamellar = 0 min). inclusions on grain boundaries. Our XRD, SEM, and EDX results contours. Pure Zr and Hf are still present in the bulk sample. suggest that these phases are Hf- and Zr-based solid solutions with interlayers of Ta- and Nb-based UponsolidUpon solutions. SPS SPS at 1500at Similar 1500◦C, °C, structures the the crystallinity crystallinity were observ of the theed bcc bccin phasethephase HEAs grew, grew, prepared although although by the arc thediffraction remelting diﬀraction peaks[38]. peaks from from Zr andZr and twoThe two solid final solid solutions product solutions sintered based based onat on 1700 Zr Zr and and°C was HfHf ( (foundhcphcp with to containaa == 3.245,3.245, two c =5.140c solid=5.140 solutionsand and bcc withbcc basedwith a = on3.455)a bcc= 3.455) andare are still presentstillhcp presentphases. in the inAs the formthe form formation of of lamellar lamellar of bcc inclusions inclusions solid solution on grain grain begins boundaries. boundaries. at 1300 °C, Our 10 Our XRD,min XRD, holding SEM, SEM, and at andEDX1300 EDXresults°C was results suggestsuggestapplied. that that these these phases phases are are Hf- Hf- and and Zr-basedZr-based solid solid solutions solutions with with interlay interlayersers of Ta- of and Ta- Nb-based and Nb-based solidsolid solutions. Figuresolutions. 3 Similar presents Similar structures thestructures XRD data were were and observedobserv SEM edimages in in the the for HEAs HEAs the bulk prepared prepared HfTaTiNbZr by byarc arcremelting specimens remelting [38]. sintered [38]. at 1300The °C forfinal 10 product min and sintered without at dwell 1700 °Ctime was (τ found= 0 min). to containAfter holding two solid for solutions 10 min, the based crystallinity on bcc and of The ﬁnal product sintered at 1700 ◦C was found to contain two solid solutions based on bcc and hcp hcpthe phases.bcc and Ashcp the phases formation increased: of bcc the solid XRD solution peaks become begins atna 1300rrower. °C, At 10 themin same holding time, at the 1300 peak °C ofwas Zr phases. As the formation of bcc solid solution begins at 1300 ◦C, 10 min holding at 1300 ◦C was applied. applied.disappeared. FigureFigure3 presents 3 presents the the XRD XRD data data and and SEM SEM images for for the the bulk bulk HfTaTiNbZr HfTaTiNbZr specimens specimens sintered sintered at at 1300 1300◦C for °C 10 for min 10 min and and without without dwell dwell time time (τ= (τ0 = min). 0 min). After After holding holding for for 10 10 min, min, the the crystallinity crystallinity of of the bcc and hcpthe phasesbcc and increased:hcp phases theincreased: XRD peaks the XRD become peaks narrower. become na Atrrower. the same At the time, same the time, peak the of peak Zr disappeared. of Zr disappeared.

Figure 3. XRD patterns and SEM images of bulk HfTaTiNbZr specimens sintered at 1300 °C for τ = 10 (a,b) and 0 min (a,c).

Figure 3. XRD patterns and SEM images of bulk HfTaTiNbZr specimens sintered at 1300 °C for τ = 10 Figure 3. XRD patterns and SEM images of bulk HfTaTiNbZr specimens sintered at 1300 ◦C for τ = 10

(a,b)( anda,b) and 0 min 0 min (a,c ().a,c).

Metals 2020, 10, 1268 6 of 15

Metals 2020, 10, x FOR PEER REVIEW 6 of 16 As follows from the data in Figures3 and4, the material obtained by SPS at 1300 ◦C for 10 min looks moreAs densefollows (Figure from the3b) data and in homogeneous Figures 3 and 4, (Figure the material4) compared obtained withby SPS that at 1300 formed °C for at 10τ min= 0 min (Figurelooks3c). more The particlesdense (Figure of Hf, 3b) Ta, and Ti, homogeneous Nb, and Zr are (Figure seen to4) becompared uniformly with distributed that formed over at τ the= 0 material,min although(Figure the 3c). traces The ofparticles Hf-rich of regions Hf, Ta, areTi, stillNb, presentand Zr are on seen grain to boundaries. be uniformly distributed over the material, although the traces of Hf-rich regions are still present on grain boundaries.

FigureFigure 4. Representative 4. Representative compositional compositional EDS mappingEDS mapping results results for the for HfTaTiNbZr the HfTaTiNbZr sample consolidatedsample at 1300consolidated◦C(τ = 10 at min). 1300 °C (τ = 10 min).

As isAs known, is known, the the mechanical mechanical alloying alloying (MA)(MA) processprocess takes takes advantage advantage of extended of extended solid solid solubility solubility eveneven in immiscible in immiscible systems systems [26]. In[26]. order In toorder improve to im theprove solubility the solubility of solids of in solids HfTaTiNbZr in HfTaTiNbZr composition, we performedcomposition, MA we in performed two diﬀerent MA in modes, two different LEBM modes, and HEBM. LEBM and HEBM.

3.2. Combined3.2. Combined Use Use of LEBM of LEBM and and SPS SPS DuringDuring LEBM LEBM (up to(up 10 to h), 10 the h), structure the structure of milled of milled HfTaTiNbZr HfTaTiNbZr powder powder blend blend underwent underwent gradual changes,Metalsgradual as2020 illustratedchanges,, 10, x FOR asPEER inillustrated FigureREVIEW5 .in Figure 5. 7 of 16

Figure 5. XRD patterns of HfTaTiNbZr powder mixtures taken after low-energy ball milling (LEBM) for diFigureﬀerent 5.t XRD(indicated patterns on of theHfTaTiNbZr right). powder mixtures taken after low-energy ball milling (LEBM) for different t (indicated on the right).

The initial HfTaTiNbZr powder blend exhibit expectedly strong and narrow Bragg peaks due to the crystalline structure of the elements. After 1 h of LEBM, the diffractions peaks of alloying elements broadened and decreased in their intensity. A further increase in milling time t up to 5 and 10 h leads to ongoing line broadening and partial overlapping of the main diffraction peaks from alloying elements. The formation of background around the (110), (200), (211), (220), and (310) peaks of the bcc phase is also observed. Most likely, this is caused by transformation of initially microcrystalline powders to nanocrystalline solid solutions and partial disordering of crystal structure upon mechanical damage. Figures 6 and 7 show the representative compositional EDX mapping results for the HfTaTiNbZr sample ball milled for 5 and 10 h, respectively. In 5 h of LEBM, Hf, Ta, Ti, Nb, and Zr form separate agglomerates of lаyеrеd structure (Figure 6), which is typical fоr ductile metal blends processed in а planetary ball mill [39]. Further mechanical treatment (Figure 7) leads to a decrease in the average layer thickness down to a nano scale. According to the EDX data, the non-uniform distribution of alloying elements is retained even after 10 h of LEBM.

Metals 2020, 10, 1268 7 of 15

The initial HfTaTiNbZr powder blend exhibit expectedly strong and narrow Bragg peaks due to the crystalline structure of the elements. After 1 h of LEBM, the diﬀractions peaks of alloying elements broadened and decreased in their intensity. A further increase in milling time t up to 5 and 10 h leads to ongoing line broadening and partial overlapping of the main diﬀraction peaks from alloying elements. The formation of background around the (110), (200), (211), (220), and (310) peaks of the bcc phase is also observed. Most likely, this is caused by transformation of initially microcrystalline powders to nanocrystalline solid solutions and partial disordering of crystal structure upon mechanical damage. Figures6 and7 show the representative compositional EDX mapping results for the HfTaTiNbZr sample ball milled for 5 and 10 h, respectively. In 5 h of LEBM, Hf, Ta, Ti, Nb, and Zr form separate agglomerates of layered structure (Figure6), which is typical for ductile metal blends processed in a planetary ball mill [39]. Further mechanical treatment (Figure7) leads to a decrease in the average layer thickness down to a nano scale. According to the EDX data, the non-uniform distribution of alloying elements is retained even after 10 h of LEBM. Metals 20202020,, 1010,, xx FORFOR PEERPEER REVIEWREVIEW 88 ofof 1616

Figure 6.FigureRepresentative 6. Representative compositional compositional EDS EDS mapping mapping results results for for the the HfTaTiNbZr HfTaTiNbZr sample sample ball milled ball milled for 5 h. forfor 55 h.h.

Figure 7.FigureRepresentative 7. Representative compositional compositional EDS EDS mapping mapping results results for for the the HfTaTiNbZr HfTaTiNbZr sample sample ball milled ball milled for 10 h.forfor 1010 h.h.

Figure 8 presents the XRD patterns and SEM images of LEBM-processed (10 h) HfTaTiNbZr specimensspecimens sinteredsintered atat 13001300 °C°C withwith dwelldwell timetime τ == 1010 minmin (a,b)(a,b) andand τ = 0 (a). Just as in the case of the bulk HfTaTiNbZr materials derived from non-milled powder blends, consolidation of LEBM-processed nanocrystalline powder blend at 1300 °C°C forfor 1010 minmin leadsleads toto anan increaseincrease inin thethe solubilitysolubility ofof alloyingalloying metalsmetals andand inin thethe crystallinitycrystallinity ofof bcc andand hcp phases.phases.

Metals 2020, 10, 1268 8 of 15

Figure8 presents the XRD patterns and SEM images of LEBM-processed (10 h) HfTaTiNbZr specimens sintered at 1300 ◦C with dwell time τ = 10 min (a,b) and τ = 0 (a). Just as in the case of the bulk HfTaTiNbZr materials derived from non-milled powder blends, consolidation of LEBM-processed nanocrystalline powder blend at 1300 ◦C for 10 min leads to an increase in the solubility of alloying metals and in the crystallinity of bcc and hcp phases. Metals 2020, 10, x FOR PEER REVIEW 9 of 16 Metals 2020, 10, x FOR PEER REVIEW 9 of 16

Figure 8. XRD patterns and SEMSEM imagesimages ofof LEBM-processedLEBM-processed (10 (10 h) h) HfTaTiNbZr HfTaTiNbZr specimens specimens sintered sintered at Figure 8. XRD patterns and SEM images of LEBM-processed (10 h) HfTaTiNbZr specimens sintered at1300 1300◦C °C with with dwell dwell time timeτ = τ 10= 10 min min (a ,(ba,b) and) andτ =τ 0= min0 min (a ().a). at 1300 °C with dwell time τ = 10 min (a,b) and τ = 0 min (a). The EDX data suggest that the elemental distributi distributionon in in the the bulk HfTaTiNbZr HfTaTiNbZr material is most homogeneous,The EDX databut contains suggest some that the small elemental Hf- and distributi Zr-rich areason in (cf.the Figure bulk HfTaTiNbZr9 9).). material is most homogeneous, but contains some small Hf- and Zr-rich areas (cf. Figure 9).

Figure 9. Representative compositional EDX mapping results for the HfTaTiNbZr sample Figure 9. Representative compositional EDX mapping results for the HfTaTiNbZr sample consolidated consolidatedFigure 9. Representative at 1300 °C for 10compositional min after 10 h ofEDX LEBM. mapping results for the HfTaTiNbZr sample at 1300 C for 10 min after 10 h of LEBM. consolidated◦ at 1300 °C for 10 min after 10 h of LEBM. On the basis of the above data, we can conclude that the use of LEBM has little or no positive On the basis of the above data, we can conclude that the use of LEBM has little or no positive influenceOn the on thebasis structure of the above of resultant data, we HfTaTiNbZr can conclude alloys. that the use of LEBM has little or no positive inﬂuence on the structure of resultant HfTaTiNbZr alloys. influence on the structure of resultant HfTaTiNbZr alloys. 3.3. Combined Use of HEBM and SPS 3.3. Combined Use of HEBM and SPS HEBM processing of HfTaTiNbZr powder blends was found to bring about significant structuralHEBM changes processing in their of comp HfTaTiNbZrosition. Figure powder 10 showsblends the was diffraction found patternsto bring takenabout after significant HEBM forstructural different changes t (indicated). in their After comp 40osition. min of Figure HEBM, 10 showsthe elemental the diffraction Bragg peaks patterns disappeared taken after and HEBM the formationfor different of solidt (indicated). solutions After with 40a hcp min (Hf-based) of HEBM, and the bcc elemental (Ta-based) Bragg structure peaks began.disappeared Moreover, and wethe observedformation the of diffractionsolid solutions peak with from a hcpthe (Hf-based)Fe that had and been bcc milled (Ta-based) from structuresteel drums began. and Moreover,balls. Longer we HEBMobserved (t >the 60 diffractionmin) leads peakto the from merging the Fe of that the hapeaksd been from milled Hf-based from (steelhcp) anddrums Ta-based and balls. (bcc Longer) solid solutionsHEBM (t and> 60 the min) formation leads to of the a commonmerging asymmetricof the peaks peak from in Hf-based the range ( hcp2θ )= and35–51°. Ta-based (bcc) solid solutionsSignals and from the formationbcc HfTaTiNbZr of a common solid solution asymmetric (Figure peak 10 in) were the range detected 2θ = after35–51°. 90 min of HEBM. Note Signalsthat the fromsignal bcc from HfTaTiNbZr Fe decreased solid slowly solution with (Figure increasing 10) twere, evidently detected to everafter growing 90 min ofthickness HEBM. Note that the signal from Fe decreased slowly with increasing t, evidently to ever growing thickness

Metals 2020, 10, 1268 9 of 15

3.3. Combined Use of HEBM and SPS HEBM processing of HfTaTiNbZr powder blends was found to bring about significant structural changes in their composition. Figure 10 shows the diffraction patterns taken after HEBM for different t (indicated). After 40 min of HEBM, the elemental Bragg peaks disappeared and the formation of solidMetals solutions2020, 10, x FOR with PEER a hcp REVIEW(Hf-based) and bcc (Ta-based) structure began. Moreover, we observed10 of 16 the diffraction peak from the Fe that had been milled from steel drums and balls. Longer HEBM (oft > milled60 min material) leads to on the the merging surface ofof thedrums peaks and from balls Hf-based during HEBM. (hcp) and In Ta-basedaddition, (somebcc) solid small solutions amount andof iron the could formation partially of a commondissolve in asymmetric the bcc HfTaTiNbZr peak in the alloy. range 2θ = 35–51◦.

Figure 10. XRD patterns of HfTaTiNbZr powder mixtures taken after high-energy ball milling (HEBM) forFigure diﬀerent 10. XRDt (indicated patterns on theof HfTaTiNbZr right). The spectra powder are mixt shiftedures along taken the aftery axis high-energy for the sake ball of clarity. milling (HEBM) for different t (indicated on the right). The spectra are shifted along the y axis for the sake of Signalsclarity. from bcc HfTaTiNbZr solid solution (Figure 10) were detected after 90 min of HEBM. Note that the signal from Fe decreased slowly with increasing t, evidently to ever growing thickness of milledThe material XRD pattern on the surfaceof RHEA of drumsalloy SPS and consolidated balls during HEBM.at 1300 In°C, addition, τ = 10 min some (Figure small amount11) clearly of ironshows could that partially the resultant dissolve material in the bccis comprisedHfTaTiNbZr of alloy.three phases: (a) bcc HfZrTiTaNb solid solution (mainThe phase); XRD pattern(b) hcpof Hf- RHEA and alloyZr-based SPS consolidatedsolid solution; at and 1300 (c)◦C, a τminor= 10 min amount (Figure of 11the) clearlyMe2Fe showsphase, thatwhere the Ме resultant = Hf, Zr. material is comprised of three phases: (a) bcc HfZrTiTaNb solid solution (main phase); (b) hcp Hf- and Zr-based solid solution; and (c) a minor amount of the Me2Fe phase, where Me = Hf, Zr.

MetalsMetals 20202020,, 1010,, 1268x FOR PEER REVIEW 1011 ofof 1516

Figure 11. XRD patterns of HfTaTiNbZr powder after HEBM for t = 90 min (lower curve) and of bulk

sampleFigure 11. sintered XRD patterns at 1300 ◦ Cof forHfTaTiNbZrτ = 10 min. powder SPS, spark after plasma HEBM sintering.for t = 90 min (lower curve) and of bulk sample sintered at 1300 °C for τ = 10 min. SPS, spark plasma sintering. The presence of (Hf, Zr)2Fe may aﬀect the physicochemical parameters of HEAs such as electronic behavior,The electricalpresence conductivity, of (Hf, Zr)2 magneticFe may properties,affect the andphysicochemical hydrogen sorption parameters/desorption of HEAs kinetics such [40,41 as]. Ourelectronic multicomponent behavior, electrical composite conductivity, can be recommended magnetic notproperties, only for structuraland hydrogen applications, sorption/desorption but also as a materialkinetics for[40,41]. hydrogen Our storagemulticomponent [42]. composite can be recommended not only for structural applications,Figure 12 but shows also as the a SEMmaterial image for hydrogen and EDX mappingstorage [42]. results for the bulk HfTaTiNbZr sample (t = 90Figure min, SPS12 shows at 1300 the◦C, SEMτ = image10 min), and illustrating EDX mapping uniform results distribution for the bulk of the HfTaTiNbZr principal elements. sample (t = 90 min,Therefore, SPS at 1300 it turns °С, out τ = that10 min), SPS-produced illustrating alloys uniform always dist containribution the of inclusions the principal of Hf- elements. and Zr-based solid solutions. In the absence of MA and in the case of LEBM, this can happen by the following scheme. During warmup, hcp and bcc metals form two solutions. At some temperature T, the hcp lattice undergoes transformation into a bcc one, which is followed by (a) mutual inter-dissolution of Ta- and Nb-based alloys and (b) formation of Ta/Nb-based and Hf/Zr-based HEAs. In the case of τ = 0, the inter-dissolution remains incomplete and the resultant material exhibits the presence of inclusions and respective diﬀraction lines (cf. Figures3 and4c). Such a scheme is also supported by a more uniform structure of materials sintered at higher T or longer dwell time. In the case of HEBM, the resultant materials turn uniform. Thermodynamics predicts that our alloy is non-equilibrium at temperatures below 800 ◦C. In view of this, during heating, it decomposes into the following three phases: one bcc Ta/Nb phase and two Hf/Zr-based phases. A high amount of structural defects acquired during the MA non-equilibrium nature of the phases favors the formation of numerous Hf-/Zr-based inclusions. The LEBM or HEBM powders sintered below 1300 ◦C showed the presence of rounded inclusions because of incomplete dissolution during the formation of the RHEA phase. In the case of non-milled powder blends, the formation of needle-like and lamellar inclusions on the grain boundaries was observed.

Metals 2020, 10, 1268 11 of 15

Metals 2020, 10, x FOR PEER REVIEW 12 of 16

Metals 2020, 10, x FOR PEER REVIEW 13 of 16

350 °C and 700 °C for LEBM powders and around 350 °C and 620 °C for HEBM powders, Figurerespectively. 12.FigureRepresentative 12. Representative compositional compositional EDX EDX mapping mapping results results for the forHfTaTiNbZr the HfTaTiNbZr sample prepared sample prepared On the basis of the XRD results for LEBM (Figure 13b) and HEBMτ (Figure 13c) powders, we by the combinedby the combined use ofuse HEBM of HEBM (t (t= = 9090 min) min) and and SPS SPS consolidation consolidation at 1300 at°C 1300( = 10 min).C(τ = 10 min). may conclude that the thermal effects at 320 °C and 350 °C have nothing ◦in common with phase transformations,Therefore, it turns but canout bethat associated SPS-produced with partial alloys ox alwaysidation containby traces the of oxygeninclusions in the of calorimeter.Hf- and TheZr-based DSCExothermic curves solid solutions.peaks for initialat 700 In the (curve°C forabsence LEBM 1), LEBMof powders MA and (curve an ind the at 2), 620case and °C of HEBMforLEBM, HEBM this (curve powders can happen 3) HfZrTiTaNbcorrespond by the to powders are presentedfollowingpartial in scheme.crystallization. Figure During13a. DSCThe warmup, XRD curves peaks hcp 1and of and bcc bcc phase 2metals exhibit become form two two narrower low-exothermic solution ands. At stronger. some peakstemperature around T, 350 ◦C and the hcp lattice undergoes transformation into a bcc one, which is followed by (a) mutual 700 ◦C for LEBMNo doubt, powders the understanding and around of 350 the ◦phases’C and formatio 620 ◦Cn forand HEBMtheir stability powders, as a function respectively. of time in inter-dissolutionRHEA HfTaTiNbZr of Ta- materials and Nb-based deserv alloyse further an ddetailed (b) formation investigation. of Ta/Nb-based and Hf/Zr-based HEAs. In the case of τ = 0, the inter-dissolution remains incomplete and the resultant material exhibits the presence of inclusions and respective diffraction lines (cf. Figures 3 and 4с). Such a scheme is also supported by a more uniform structure of materials sintered at higher T or longer dwell time. In the case of HEBM, the resultant materials turn uniform. Thermodynamics predicts that our alloy is non-equilibrium at temperatures below 800 °С. In view of this, during heating, it decomposes into the following three phases: one bcc Ta/Nb phase and two Hf/Zr-based phases. A high amount of structural defects acquired during the MA non-equilibrium nature of the phases favors the formation of numerous Hf-/Zr-based inclusions. The LEBM or HEBM powders sintered below 1300 °C showed the presence of rounded inclusions because of incomplete dissolution during the formation of the RHEA phase. In the case of non-milled powder blends, the formation of needle-like and lamellar inclusions on the grain boundaries was observed. The DSC curves for initial (curve 1), LEBM (curve 2), and HEBM (curve 3) HfZrTiTaNb powders are presented in Figure 13a. DSC curves 1 and 2 exhibit two low-exothermic peaks around

Figure 13. Figure(a) DSC 13. ( curvesa) DSC curves for HfZrTiTaNb for HfZrTiTaNb powders: powders: (1) init initial,ial, (2) (2)after after 90 min 90 of min HEBM, of HEBM,and (3) after and (3) after 10 h of LEBM;10 h (ofb )LEBM; XRD ( patternsb) XRD patterns of LEBM-processed of LEBM-processed HfTaTiNbZr HfTaTiNbZr powder powder mixtures mixtures at 350 at °C 350 and◦C 620 and 620 ◦C; and (c) XRD°C; patternsand (c) XRD of patterns HEBM-processed of HEBM-processed HfTaTiNbZr HfTaTiNbZr powder powder mixtures mixtures at at 350 350 °C◦ andC and 700 °C. 700 ◦C.

The Vickers hardness of RHEA materials prepared by HEBM (t = 90 min) and SPS consolidation (1300 °C, τ =10 min) turned markedly higher (10.7 GPa) than that of SPS-produced from unprocessed elemental powders (6.4 GPa) or from LEBM (t = 10 h) powders (5.8 GPa). Moreover, this is the highest value of Vickers hardness of all HEAs listed in [43] and three times higher than that for the HfTaTiNbZr alloy prepared by acr melting [24]. The density of our HfTaTiNbZr alloy is 9.87 ± 0.18 g/cm³ (98.7%). Superhigh hardness of our alloy can be associated with the strengthening of (a) solid solution caused by strong lattice distortion [10] and (b) grain boundaries with ultrafine grains. A similar increase in mechanical properties was reported for SPSed WNbMoTaV [44].

Metals 2020, 10, 1268 12 of 15

On the basis of the XRD results for LEBM (Figure 13b) and HEBM (Figure 13c) powders, we may conclude that the thermal eﬀects at 320 ◦C and 350 ◦C have nothing in common with phase transformations, but can be associated with partial oxidation by traces of oxygen in the calorimeter. Exothermic peaks at 700 ◦C for LEBM powders and at 620 ◦C for HEBM powders correspond to partial crystallization. The XRD peaks of bcc phase become narrower and stronger. No doubt, the understanding of the phases’ formation and their stability as a function of time in RHEA HfTaTiNbZr materials deserve further detailed investigation. The Vickers hardness of RHEA materials prepared by HEBM (t = 90 min) and SPS consolidation (1300 ◦C, τ =10 min) turned markedly higher (10.7 GPa) than that of SPS-produced from unprocessed elemental powders (6.4 GPa) or from LEBM (t = 10 h) powders (5.8 GPa). Moreover, this is the highest value of Vickers hardness of all HEAs listed in [43] and three times higher than that for the HfTaTiNbZr alloy prepared by acr melting [24]. The density of our HfTaTiNbZr alloy is 9.87 0.18 g/cm3 (98.7%). ± Superhigh hardness of our alloy can be associated with the strengthening of (a) solid solution caused by strong lattice distortion [10] and (b) grain boundaries with ultraﬁne grains. A similar increase in mechanical properties was reported for SPSed WNbMoTaV [44].

4. Conclusions (1) Powder of refractory bcc HfTaTiNbZr-based high-entropy alloy (RHEA) was successfully produced by short-term (90 min) HEBM. The evolution in the structure of initial HfTaTiNbZr powder mixture during HEBM was elucidated by XRD. All Bragg peaks of principal elements, intermediate solid solutions with a hcp (Hf-based) and bcc (Ta-based) structure, vanished in 90 min of HEBM, forming HEA solid solution with a bcc structure. (2) Consolidation of HEBM-produced RHEA powder by SPS at 1300 ◦C for 10 min results in the formation of bulk material with an ultra-high Vickers hardness (10.7 GPa) and a density of 9.87 0.18 g/cm3 (98.7%). ± (3) On the basis of the XRD and SEM/EDX results, HfTaTiNbZr-based RHEA alloy was found to consist of bcc HfZrTiTaNb as a main phase, hcp Hf/Zr-based solid solution, and Me2Fe phases (Me = Hf, Zr) as minor admixtures. Principal elements of the HEA phase were uniformly distributed over the bulk of HfTaTiNbZr alloy. (4) The HfTaTiNbZr alloys synthesized in the absence of ball milling and in the case of low-energy ball milling (LEBM, 10 h) consisted of bcc HEA and Hf/Zr-rich hcp solid solution; in this case, the Vickers hardness of such alloys was found to have a value of 6.4 GPa and 5.8 GPa, respectively. (5) An optimal combination of HEBM and SPS processes can be recommended as a facile route for fabrication of equiatomic refractory HfTaTiNbZr-based high-entropy powders and bulk materials with the best structural/elemental homogeneity and promising mechanical properties.

Author Contributions: A.S., S.B., and D.M. conceived of and designed the experiments. A.S., S.B., S.V., and K.K. performed the experiments. N.S., K.K., Y.S., and D.M. analyzed the data and wrote the paper. All authors have read and agreed to the published version of the manuscript. Funding: This research was funded by the Russian Science Foundation, Grant Number 18-79-10215. Acknowledgments: This work was conducted with the financial support of the Russian Science Foundation (Grant Number 18-79-10215). Conflicts of Interest: The authors declare no conflict of interest.

References

1. Yeh, J.-W.; Chen, S.-K.; Lin, S.-J.; Gan, J.-Y.; Chin, T.-S.; Shun, T.-T.; Tsau, C.-H.; Chang, S.-Y. Nanostructured High-Entropy Alloys with Multiple Principal Elements: Novel Alloy Design Concepts and Outcomes. Adv. Eng. Mater. 2004, 6, 299–303. [CrossRef] 2. Cantor, B.; Chang, I.T.H.; Knight, P.; Vincent, A.J.B. Microstructural development in equiatomic multicomponent alloys. Mater. Sci. Eng. A 2004, 375–377, 213–218. [CrossRef] Metals 2020, 10, 1268 13 of 15

3. Gludovatz, B.; Hohenwarter, A.; Catoor, D.; Chang, E.H.; George, E.P.; Ritchie, R.O. A fracture-resistant high-entropy alloy for cryogenic applications. Science 2014, 345, 1153–1158. [CrossRef][PubMed] 4. Zhang, Z.; Mao, M.M.; Wang, J.; Gludovatz, B.; Zhang, Z.; Mao, S.X.; George, E.P.; Yu, Q.; Ritchie, R.O. Nanoscale origins of the damage tolerance of the high-entropy alloy CrMnFeCoNi. Nat. Commun. 2015, 6, 10143. [CrossRef][PubMed] 5. Gali, A.; George, E.P. Tensile properties of high- and medium-entropy alloys. Intermetallics 2013, 39, 74–78. [CrossRef] 6. Otto, F.; Dlouhý, A.; Somsen, C.; Bei, H.; Eggeler, G.; George, E.P. The influences of temperature and microstructure on the tensile properties of a CoCrFeMnNi high-entropy alloy. Acta Mater. 2013, 61, 5743–5755. [CrossRef] 7. Kilmametov, A.; Kulagin, R.; Mazilkin, A.; Seils, S.; Boll, T.; Heilmaier, M.; Hahn, H. High-pressure torsion driven mechanical alloying of CoCrFeMnNi high entropy alloy. Scr. Mater. 2019, 158, 29–33. [CrossRef] 8. Zhu, C.; Lu, Z.P.; Nieh, T.G. Incipient plasticity and dislocation nucleation of FeCoCrNiMn high-entropy alloy. Acta Mater. 2013, 61, 2993–3001. [CrossRef] 9. Sun, S.J.; Tian, Y.Z.; Lin, H.R.; Dong, X.G.; Wang, Y.H.; Zhang, Z.J.; Zhang, Z.F. Enhanced strength and ductility of bulk CoCrFeMnNi high entropy alloy having fully recrystallized ultrafine-grained structure. Mater. Des. 2017, 133, 122–127. [CrossRef] 10. Miracle, D.B.; Senkov, O.N. A critical review of high entropy alloys and related concepts. Acta Mater. 2017, 122, 448–511. [CrossRef] 11. Chang, X.; Zeng, M.; Liu, K.; Fu, L. Phase Engineering of High-Entropy Alloys. Adv. Mater. 2020, 32, 1907226. [CrossRef][PubMed] 12. GUO, S.; LIU, C.T. Phase stability in high entropy alloys: Formation of solid-solution phase or amorphous phase. Prog. Nat. Sci. Mater. Int. 2011, 21, 433–446. [CrossRef] 13. Guo, S.; Ng, C.; Lu, J.; Liu, C.T. Effect of valence electron concentration on stability of fcc or bcc phase in high entropy alloys. J. Appl. Phys. 2011, 109, 103505. [CrossRef] 14. Mishra, R.K.; Shahi, R.R. Magnetic Characteristics of High Entropy Alloys. In Magnetism and Magnetic Materials; InTech: London, UK, 2018. 15. Zhang, K.; Fu, Z. Effects of annealing treatment on phase composition and microstructure of CoCrFeNiTiAlx high-entropy alloys. Intermetallics 2012, 22, 24–32. [CrossRef] 16. Zherebtsov, S.; Stepanov, N.; Ivanisenko, Y.; Shaysultanov, D.; Yurchenko, N.; Klimova, M.; Salishchev, G. Evolution of Microstructure and Mechanical Properties of a CoCrFeMnNi High-Entropy Alloy during High-Pressure Torsion at Room and Cryogenic Temperatures. Metals (Basel) 2018, 8, 123. [CrossRef] 17. Rogachev, A.S.; Vadchenko, S.G.; Kochetov, N.A.; Rouvimov, S.; Kovalev, D.Y.; Shchukin, A.S.; Moskovskikh, D.O.; Nepapushev, A.A.; Mukasyan, A.S. Structure and properties of equiatomic CoCrFeNiMn alloy fabricated by high-energy ball milling and spark plasma sintering. J. Alloys Compd. 2019, 805, 1237–1245. [CrossRef] 18. Wang, H.; Gao, T.; Niu, J.; Shi, P.; Xu, J.; Wang, Y. Microstructure, thermal properties, and corrosion behaviors of FeSiBAlNi alloy fabricated by mechanical alloying and spark plasma sintering. Int. J. Miner. Metall. Mater. 2016, 23, 77–82. [CrossRef] 19. Murty, B.S.; Yeh, J.-W.; Ranganathan, S.; Bhattacharjee, P.P. High-Entropy Alloys; Elsevier: London, UK, 2019; ISBN 9780128160671. 20. Gao, M.C.; Yeh, J.-W.; Liaw, P.K.; Zhang, Y. High-Entropy Alloys Fundamentals and Applications; Gao, M.C., Yeh, J.-W., Liaw, P.K., Zhang, Y., Eds.; Springer International Publishing: Cham, Swizerland, 2016; ISBN 978-3-319-27011-1. 21. Syed Ghazi, S.; Ravi, K.R. Phase-evolution in high entropy alloys: Role of synthesis route. Intermetallics 2016, 73, 40–42. [CrossRef] 22. Panina, E.; Yurchenko, N.; Zherebtsov, S.; Stepanov, N.; Salishchev, G.; Ventzke, V.; Dinse, R.; Kashaev, N. Laser Beam Welding of a Low Density Refractory High Entropy Alloy. Metals (Basel) 2019, 9, 1351. [CrossRef] 23. Yurchenko, N.; Panina, E.; Zherebtsov, S.; Salishchev, G.; Stepanov, N. Oxidation Behavior of Refractory AlNbTiVZr0.25 High-Entropy Alloy. Materials (Basel) 2018, 11, 2526. [CrossRef] 24. Senkov, O.N.; Scott, J.M.; Senkova, S.V.; Miracle, D.B.; Woodward, C.F. Microstructure and room temperature properties of a high-entropy TaNbHfZrTi alloy. J. Alloys Compd. 2011, 509, 6043–6048. [CrossRef] Metals 2020, 10, 1268 14 of 15

25. Liu, Q.; Wang, G.; Sui, X.; Liu, Y.; Li, X.; Yang, J. Microstructure and mechanical properties of ultra-fine grained MoNbTaTiV refractory high-entropy alloy fabricated by spark plasma sintering. J. Mater. Sci. Technol. 2019, 35, 2600–2607. [CrossRef] 26. Shkodich, N.F.; Rogachev, A.S.; Vadchenko, S.G.; Moskovskikh, D.O.; Sachkova, N.V.; Rouvimov, S.; Mukasyan, A.S. Bulk Cu-Cr nanocomposites by high-energy ball milling and spark plasma sintering. J. Alloys Compd. 2014, 617, 39–46. [CrossRef] 27. Murty, B.S.; Ranganathan, S. Novel materials synthesis by mechanical alloying/milling. Int. Mater. Rev. 1998, 43, 101–141. [CrossRef] 28. Shkodich, N.F.; Vadchenko, S.G.; Nepapushev, A.A.; Kovalev, D.Y.; Kovalev, I.D.; Ruvimov, S.; Rogachev, A.S.; Mukasyan, A.S. Crystallization of amorphous Cu50Ti50 alloy prepared by high-energy ball milling. J. Alloys Compd. 2018, 741, 575–579. [CrossRef] 29. Rogachev, A.S. Mechanical activation of heterogeneous exothermic reactions in powder mixtures. Russ. Chem. Rev. 2019, 88, 875–900. [CrossRef] 30. Rogachev, A.S.; Kuskov, K.V.; Moskovskikh, D.O.; Usenko, A.A.; Orlov, A.O.; Shkodich, N.F.; Alymov, M.I.; Mukasyan, A.S. Effect of mechanical activation on thermal and electrical conductivity of sintered Cu, Cr, and Cu/Cr composite powders. Dokl. Phys. 2016, 61.[CrossRef] 31. Varalakshmi, S.; Kamaraj, M.; Murty, B.S. Synthesis and characterization of nanocrystalline AlFeTiCrZnCu high entropy solid solution by mechanical alloying. J. Alloys Compd. 2008, 460, 253–257. [CrossRef] 32. Cavaliere, P.; Sadeghi, B.; Shabani, A. Spark Plasma Sintering: Process Fundamentals. In Spark Plasma Sintering of Materials: Advances in Processing and Applications; Springer Nature: Basel, Switzerland, 2019; ISBN 9783030053277. 33. Hu, Z.-Y.; Zhang, Z.-H.; Cheng, X.-W.; Wang, F.-C.; Zhang, Y.-F.; Li, S.-L. A review of multi-physical fields induced phenomena and effects in spark plasma sintering: Fundamentals and applications. Mater. Des. 2020, 108662. [CrossRef] 34. Dudina, D.V.; Bokhonov, B.B.; Olevsky, E.A. Fabrication of porous materials by spark plasma sintering: A review. Materials (Basel) 2019, 12, 541. [CrossRef] 35. Abedi, M.; Moskovskikh, D.O.; Rogachev, A.S.; Mukasyan, A.S. Spark Plasma Sintering of Titanium Spherical Particles. Metall. Mater. Trans. B Process Metall. Mater. Process. Sci. 2016, 47.[CrossRef] 36. Schuh, B.; Völker, B.; Todt, J.; Schell, N.; Perrière, L.; Li, J.; Couzinié, J.P.; Hohenwarter, A. Thermodynamic instability of a nanocrystalline, single-phase TiZrNbHfTa alloy and its impact on the mechanical properties. Acta Mater. 2018, 142, 201–212. [CrossRef] 37. Singh, A.K.; Subramaniam, A. On the formation of disordered solid solutions in multi-component alloys. J. Alloys Compd. 2014, 587, 113–119. [CrossRef] 38. Chen, S.Y.; Tong, Y.; Tseng, K.-K.; Yeh, J.-W.; Poplawsky, J.D.; Wen, J.G.; Gao, M.C.; Kim, G.; Chen, W.; Ren, Y.; et al. Phase transformations of HfNbTaTiZr high-entropy alloy at intermediate temperatures. Scr. Mater. 2019, 158, 50–56. [CrossRef] 39. Rogachev, A.S.; Shkodich, N.F.; Vadchenko, S.G.; Baras, F.; Kovalev, D.Y.; Rouvimov, S.; Nepapushev, A.A.; Mukasyan, A.S. Influence of the high energy ball milling on structure and reactivity of the Ni + Al powder mixture. J. Alloys Compd. 2013, 577, 600–605. [CrossRef] 40. Martinez-Amariz, A.; Peña, D.; Montes Vera, E. Metal hydrides with body-centered cubic structure: Advantages and challenges. J. Phys. Conf. Ser. 2019, 1386, 012061. [CrossRef]

41. Zavaliy, I.Y.; Riabov, A.; Yartys, V.; Wiesinger, G.; Michor, H.; Hilscher, G. (Hf,Zr)2Fe and Zr4Fe2O0.6 compounds and their hydrides: Phase equilibria, crystal structure and magnetic properties. J. Alloys Compd. 1998, 265, 6–14. [CrossRef] 42. Zadorozhnyy,V.;Sarac, B.; Berdonosova, E.; Karazehir, T.; Lassnig, A.; Gammer, C.; Zadorozhnyy,M.; Ketov,S.; Klyamkin, S.; Eckert, J. Evaluation of hydrogen storage performance of ZrTiVNiCrFe in electrochemical and gas-solid reactions. Int. J. Hydrogen Energy 2020, 45, 5347–5355. [CrossRef] Metals 2020, 10, 1268 15 of 15

43. Gorsse, S.; Nguyen, M.H.; Senkov, O.N.; Miracle, D.B. Database on the mechanical properties of high entropy alloys and complex concentrated alloys. Data Br. 2018, 21, 2664–2678. [CrossRef] 44. Kang, B.; Lee, J.; Ryu, H.J.; Hong, S.H. Ultra-high strength WNbMoTaV high-entropy alloys with ﬁne grain structure fabricated by powder metallurgical process. Mater. Sci. Eng. A 2018, 712, 616–624. [CrossRef]

© 2020 by the authors. Licensee MDPI, Basel, Switzerland. This article is an open access article distributed under the terms and conditions of the Creative Commons Attribution (CC BY) license (http://creativecommons.org/licenses/by/4.0/).