Development of a TiNbTaMoZr-Based High Entropy Alloy with Low Young´s Modulus by Mechanical Alloying Route

metals Article Development of a TiNbTaMoZr-Based High Entropy Alloy with Low Young´s Modulus by Mechanical Alloying Route Juliette Normand 1, Rocío Moriche 2, Cristina García-Garrido 3, Ranier Enrique Sepúlveda Ferrer 4and Ernesto Chicardi 4,* 1SIGMA Clermont, 63178 Aubière, France; [email protected] 2Departamento de Física de la Materia Condensada, Universidad de Sevilla, 41080 Sevilla, Spain; [email protected] 3Instituto Andaluz del Patrimonio Histórico (IAPH), 41092 Sevilla, Spain; [email protected] 4Departamento de Ingeniería y Ciencia de los Materiales y del Transporte, Universidad de Sevilla, 41092 Sevilla, Spain; [email protected] *Correspondence: [email protected]; Tel.: +34-954-482-279 Received: 14 October 2020; Accepted: 29 October 2020; Published: 1 November 2020   Abstract: In this work, an equiatomic TiNbTaMoZr-based high-entropy alloy (HEA) has been developed by a powder metallurgy route, which consists of a process of combined one-step low-temperature mechanical milling starting from the transition metals as raw materials and a subsequent pressureless sintering. In this way, the optimized synthesized specimen, after 10 h of milling time, showed two different body-centered cubic (bcc) TiNbTaMoZr alloys, which, after sintering at 1450 ◦ C, 1 h of dwell time and a heating and cooling rate of 5 ◦ C min −1 , it remained formed as two bcc TiNbTaMoZr-based HEAs. This material, with micrometric and equiaxed particles, and with homogeneously distributed phases, presented a Young’s modulus that was significantly higher (5.8 GPa) and lower (62.1 GPa) than that of the usual commercially pure (cp) Ti and Ti6Al4V alloy used for bone-replacement implants. It also presented similar values to those of the HEAs developed for the same purpose. These interesting properties would enable this TiNbTaMoZr-based HEA to be used as a potential biomaterial for bulk or porous bone implants with high hardness and low Young´s modulus, thereby preventing the appearance of stress-shielding phenomena. Keywords: high-entropy alloys; titanium; biomaterials; bone tissue engineering; mechanical alloying; Young’s modulus 1. Introduction High-entropy alloys (HEAs) form a new group of metallic materials that has been under investigation over the past few decades. Unlike conventional alloys based on a single main element, HEAs are composed of at least five elements in near-equiatomic proportions [ 1 ]. In 2013, Yeh [ 2 ] proposed HEAs as those alloys with a configurational entropy of mixing ( ∆ S mix ) above 1.5R. For an equiatomic alloy, ∆Smix is defined as: ∆Smix =Rln(n)(1) with nas the number of elements and Rthe gas constant. Therefore, in order to obtain a ∆Smix above 1.5R, an HEA would have to contain at least five elements. Metals 2020,10, 1463; doi:10.3390/met10111463 www.mdpi.com/journal/metals Metals 2020,10, 1463 2 of 15 The success of these new alloys is mainly based on their high stability, provided by the high mixing entropies ( ∆ S mix ) for solid solutions, which compensate for the positive enthalpy of mixing (∆Hmix) to obtain a negative Gibbs free energy of mixing (∆Gmix) [2]. ∆Gmix =∆Hmix −T∆Smix (2) To promote the formation of HEAs, Zhang et al. in 2008 [ 3 ] and Guo et al. in 2013 [ 4 ] also point out the importance of the atomic size difference (δ), according to Equation (3). δ=100v tn X i xi1−ri r2(3) where x i and r i are the atomic fraction and radius of the i th element and r is the average atomic radius. The solid solution was predicted to occur for approximately δ <6.4% associated with the enthalpy and entropy of mixing of − 20 kJ · mol −1≤∆ H mix ≤ 5 kJ · mol −1 and 12 J · mol −1· K −1≤∆ S mix ≤ 17.5 J · mol −1· K −1 respectively [3]. All these parameters have also been corroborated more recently by Ye et al. [5]. Therefore, as a consequence of the expected stabilization of the solid solution and microstructure, a variety of other unique properties arise from their multi-atom compositions [ 6 , 7 ]. For example, the complex and random distribution of atoms generates distortion and stress in the lattice, which lead to an increase of toughness, strength and wear resistance of the material [ 2 , 8 ]. Their superior mechanical properties compared to conventional alloys have attracted various engineering fields, such as the transportation and energy industries, for their use in structural or refractory applications. Certain literature is devoted to the study of HEAs as potential materials for these industries [8–10]. Regarding the biomedical field, these innovative HEAs could provide a good choice of material since they comply with the strict criteria for orthopaedic applications. Bone replacements must be biocompatible and must show no cytotoxicity and offer good resistance to corrosion within the body environment [ 11 ]. They should also exhibit a low Young’s modulus, which is associated with great strength, high elasticity and high wear resistance to mimic bone properties and prevent the stress-shielding effect that causes bone resorption and the loosening of implants [ 12 – 14 ]. The currently used Ti-based alloys, mainly Ti6Al4V, satisfy most of these criteria but exhibit an excessively high Young’s modulus (E =112 GPa) compared to human bones (E =3–40 GPa) [ 14 ]. Moreover, Al and V have proven to be toxic to the body [ 15 ]. Therefore, the biocompatible elements such as Ta, Nb, Mo, Zr, Mg and Sn are studied herein as Ti alloying elements for orthopaedic uses [15,16]. Furthermore, to reduce the value of Young’s modulus, one of the strategies is to develop β -Ti based alloys (body-centered cubic phase), which possess greater strength and a lower Young’s modulus than the α (hexagonal closed-packed structure) Ti-based alloys [ 17 – 19 ]. The allotropic transformation from α to β phases takes place at 882.3 ◦ C although the addition of alloying elements, known as β -stabilizers, such as Nb and Ta, can lower this temperature [ 11 , 19 ]. Since it is known that Ti, Nb and Ta possess complete mutual solubility and perfect biocompatibility, the association of these three elements has attracted interest as a basis for the development of alloys [ 20 – 22 ]. The formation of a HEA by combining this basis with two or more biocompatible metals would lead to a material with an even higher yield strength and lower Young’s modulus, thereby providing a material that satisfies the orthopaedic criteria in terms of mechanical properties, non-toxicity and corrosion resistance. The synthesis of the HEAs is essentially obtained by the arc-melting process, which can achieve sufficiently high temperatures to melt most of the metals used. This method usually leads to a heterogeneous microstructure with a major segregation of the atoms [23]. A less common processing route to synthesis the HEAs is the mechanical alloying (MA). The MA route is a solid-state high-energy ball milling process, which is a simple, reproducible, scalable and inexpensive synthesis route that would enable the HEA to be obtained in a powder form in contrast with the arc-melting process. The particles are plastically deformed, fractured and welded at high speed, which allows diffusion Metals 2020,10, 1463 3 of 15 of the different elements into the particles and promotes nanocrystallization. Better control of the structure and constitution of homogeneous materials can therefore be achieved [24]. The formation of a nominal equiatomic HEA composed of the biocompatible Ti, Nb, Ta, Mo and Zr metals has previously been completed using the arc-melting process by Wang and Xu [ 25 ], Todai et al. [ 26 ], and Nagase et al. [ 27 ]. They obtained a structure formed of two bcc alloys, with dendritic morphology and heterogeneous composition, due to the segregations of metals atoms during solidification; however, a noticeably high Young’s modulus of 133 GPa was also reported [25]. For a hypothetical equiatomic HEA composed of the five elements employed, that is, Ti, Nb, Ta, Mo and Zr, at 25 ◦ C, the ∆ S mix and ∆ H mix were calculated as 13.4 J · mol −1· K −1 and − 1.6 kJ · mol −1 respectively, thereby giving a ∆ G mix of − 5.6 kJ/mol following Equations (1)–(3). This negative result predicts an elevated stability for the alloy with a calculated atomic size difference δ equal to 4.9%. These values are in good agreement with the ranges determined by Zhang, et al. [ 3 ] (12 ≤∆Smix ≤17.5 J·mol−1·K−1 ; − 20 ≤∆ H mix ≤ 5 kJ/mol and δ <6.4%). This analysis suggests the possible formation of the expected TiNbTaMoZr HEA by MA, from a thermodynamic perspective. Therefore, the aim of this work is to study the synthesis, characterize the final microstructure and determined the mechanical properties of the TiNbTaMoZr based HEA by using the alternative mechanical alloying route, in order to ascertain the suitability of this material for its use as biomaterial for bulk of porousorthopedic bone implant developed by additive manufacturing techniques [28]. 2. Experimental Procedure 2.1. Development of the TiNbTaMoZr-Based HEA. Equiatomic powder mixtures of titanium (99% purity, <325 mesh, Strem chemicals, Newburyport, MA, USA), tantalum (99.6% purity, <325 mesh, Strem chemicals, Newburyport, MA, USA), niobium (99.8% purity, <325 mesh, Strem chemicals, Newburyport, MA, USA), molybdenum (99% purity, <325 mesh, Strem chemicals, Newburyport, MA, USA) and zirconium (99.5% purity, 50 mesh, Strem chemicals, Newburyport, MA, USA) were used as the raw materials to synthesize the TiNbTaMoZr-based HEA. The TiNbTaMoZr powder mixtures were milled at various milling times between 1 and 20 h, with a spinning rate of 600 rpm in a planetary ball mill (P7, Fritsch), by cycles of 15 min ON and 15 min OFF to prevent an excessive increase in temperature and also to prevent the welding of powders, under an inert argon gas atmosphere at 4 atm. (Argon, Linde Group Spain, Valencia, Spain). A 45 mL steel jar and 7 steel balls (Ø =15 mm and m =13.5 g) were used with a ball-to-powder ratio (BPR) equal to 10. The as-milled powders were labeled as pTNTMZxh, where pand xcorrespond to powders and the hours of milling time respectively. The pTNTMZ powders selected were compacted using a uniaxial press (2 tons) and, subsequently, a cold isostatic press (200 MPa). The green bodies were sintered in a horizontal tubular furnace (IGM1360 model No. RTH-180-50-1H, AGNI, Aechen, Germany) under an Argon flow (H2O 68 ppm and O 2 62 ppm, Linde Group Spain, Valencia, Spain) at 1450 ◦ C and 1550 ◦ C for 1 h to study the microstructure evolution and the possible transformation of the phases. The heating and cooling rates used were 5 ◦ C · min −1 . The final cylindrical specimens of 20 mm in diameter and 5 mm in height were obtained. These sintered specimens were labeled as sTNTMZ_y ◦ C, where sand ycorrespond to sintered and sintering temperature respectively. 2.2. Microstructural Characterization of the TiNbTaMoZr-Based HEA. X-ray diffraction (XRD) patterns of the as-synthesized pTNTMZxh powders and the as-sintered sTNTMZ_y ◦ CspecimenswererecordedusingaPANalyticalX’PertProinstrument(MalvernPanalytical Ltd, Malvern, UK) equipped with a Bragg-Brentano θ / θ geometry, a Cu K α radiation source (45 kV, 40 mA ), a secondary K β filter and an X’Celerator detector. The XRD patterns were obtained by scanning between 20 and 120 ◦ of the 2theta degree with 0.03 ◦ steps and a counting time of 800 s · step −1 . Metals 2020,10, 1463 4 of 15 Thefree-accessCrystallographyOpen Database(COD) wasemployed toelucidate thephases developed with the corresponding structures and space-group symmetries (SGSs). A lanthanum hexaboride, LaB6 (Standard Reference Material 660b, NIST) pattern was used to calibrate the positions of the diffraction lines. Scanning electron microscopy (SEM) images in secondary electron mode were obtained using a Hitachi S-4800 (Hitachinaka, Japan) field at a voltage of 5 kV on the specimens. The Ti, Nb, Ta, Mo, Zr and Fe (coming from the milling media) contents were measured by energy-dispersive X-ray spectrometry (EDS) point analysis using a detector coupled with SEM at a voltage of 20 kV. All SEM–EDS maps were collected under a voltage of 30 kV. High-resolution transmission electron microscopy (HRTEM) images were obtained by using an FEI Talos ™ F200S scanning/transmission electron microscope (Hillsboro, OR, USA) at a voltage of 200 kV (point resolution =0.25 nm). The micrograph analysis, lattice spacing, fast Fourier transform (FFT) and phase interpretation were performed with the Gatan Digital micrograph software (Gatan Inc., Pleasanton, CA, USA). The as-milled pTNTMZ powder sample was dispersed in ethanol and droplets of the suspension were deposited onto a PELCO®copper grid with pure carbon film. 2.3. Micromechanical Behavior of the TiNbTaMoZr-Based HEA. The micromechanical behavior of the sintered TiNbTaMoZr-based HEA was analyzed by instrumented microindentation tests. Static micromechanical tests were performed in a MTR3/50-50/NI Microtest durometer machine (Microtest, Madrid, Spain) using a Vickers indenter. The maximum load was 1 N with a load rate of 0.5 N/min and a dwell time of 5 s. The microhardness and Young’s modulus were calculated from the resulting loading–unloading curves following the method published by Oliver and Pharr [ 29 ]. The microhardness (H) of the HEA was calculated from the maximum load ( Pmax ) and the contact area (A), calculated from the depth of contact between the indenter and the sample surface (hc), in accordance with Equation (4): H=Pmax A(4) The effective elastic modulus ( Eef f ) was calculated by the slope of the unloading P–h curve ( S ) and a correction factor dependent on the indenter (β), using the Equation (5): Eef f =S β2 √π√A(5) Finally, the elastic modulus ( E ) was calculated from Eef f while taking into account the contribution of the elastic modulus (Ei) and Poisson’s ratio of the indenter (νi) following Equation (6): E=1−ν2 1 Eef f −(1−νi2) Ei (6) 3. Results and Discussion 3.1. Synthesis and Characterization of the TiNbTaMoZr-Based HEA. The X-ray diffraction (XRD) patterns obtained for the different as-synthesized pTNTMZ specimens, milled for between 1 and 20 h, are shown in Figure 1a. It should be noted that pTNTMZ2h, pTNTMZ3h, pTNTMZ5h and pTNTMZ6h have been excluded from the figure for the purpose of simplification and due to no substantial difference being found. Thus, for pTNTMZ1h, three phases could be indexed, that is, a hexagonal close-packed (hcp) structure, with P63/mmc SGS, corresponding to the elemental Zr (COD no. 9008523) and two body-centered cubic (bcc) structures with Im3m SGS. Both bcc structures can be assigned to elemental Nb or Ta for the first bcc, and to elemental Mo for the second bcc. Initially, Metals 2020,10, 1463 5 of 15 Nb and Ta could rapidly form a solid solution since they have the same crystal structure (COD no. 1539041 and 1541266 respectively), a similar lattice parameter, and equivalent covalent radii, while Mo (COD no. 9008543) have a distinguished lattice parameter. Regarding the magnification of the XRD patterns around the (110) peak of the Nb, Ta and Mo in pTNTMZ1h (Figure 1b), a displacement to a lower 2theta degree can be observed, which suggests that both bcc phases are modifying their compositions, attributable to an evolution from elemental Nb, Ta and Mo to two bcc-1 and bcc-2 alloys (dotted lines in Figure 1b). Taking into account the total absence of any peaks corresponding to the elemental Ti (P63/mmc SGS, COD no. 9008517) in the pTNTMZ1h XRD pattern, a plausible explanation to the formation of both bcc-1 and bcc-2 alloys is the incorporation of Ti atoms (with the smaller covalent radii) to both bcc structures due to the high Ti solubility in the Nb and Ta bcc structure, and to the Ti partial solubility in Mo, according to the corresponding binary phase diagrams [30]. Metals 2020, 10, x FOR PEER REVIEW 5 of 16 parameter. Regarding the magnification of the XRD patterns around the (110) peak of the Nb, Ta and Mo in pTNTMZ1h (Figure 1b), a displacement to a lower 2theta degree can be observed, which suggests that both bcc phases are modifying their compositions, attributable to an evolution from elemental Nb, Ta and Mo to two bcc-1 and bcc-2 alloys (dotted lines in Figure 1b). Taking into account the total absence of any peaks corresponding to the elemental Ti (P63/mmc SGS, COD no. 9008517) in the pTNTMZ1h XRD pattern, a plausible explanation to the formation of both bcc-1 and bcc-2 alloys is the incorporation of Ti atoms (with the smaller covalent radii) to both bcc structures due to the high Ti solubility in the Nb and Ta bcc structure, and to the Ti partial solubility in Mo, according 25 30 35 40 45 50 55 60 65 70 75 80 α -Fe (Im-3m) Zr (P63/mmc) bcc-2 (Im3m) bcc-1 (Im3m) ♣ ♦ * • * (110) (101) (002) (100) (211) (211) (200) (110) (200) (110) Intensity (a.u.) 2 theta ( º ) a) pTNTMZ20h pTNTMZ10h pTNTMZ8h pTNTMZ4h • • • • •• •• * * * ♦ ♦ • • ♦ ♦ * ♦ ♦ ♦ ♦ ♦ • • • • • pTNTMZ1h ♣ Figure 1. Cont. Metals 2020,10, 1463 6 of 15 Metals 2020, 10, x FOR PEER REVIEW 6 of 16 Figure 1. (a) XRD patterns showing the pTNTMZ specimens milled between 1 and 20 h. (b) Magnification of all XRD patterns and (c) fitting for the pTNTMZ10h in the 26–48° 2θ range and around the (110) crystallographic planes for the both two bcc-1 and bcc-2 alloys. The corresponding crystallographic planes for each phase are marked between brackets. Dotted lines are used to highlight the shift in the XRD peaks. 26 28 30 32 34 36 38 40 42 44 46 48 * (101) (002) (100) (110) (110) (110) b) pTNTMZ20h pTNTMZ10h pTNTMZ8h pTNTMZ4h • • • * * * • ♦ ♦ * ♦ ♦ ♦ • pTNTMZ1h ♣ Intensity (a.u.) 2 theta (º) α -Fe (Im-3m) Zr (P63/mmc) bcc-2 (Im3m) bcc-1 (Im3m) ♣ ♦ * • 26 28 30 32 34 36 38 40 42 44 46 48 c ) pTNTZM10h bcc-1 bcc-2 bcc-sum Intensity (a.u.) 2 theta ( º ) Figure 1. ( a ) XRD patterns showing the pTNTMZ specimens milled between 1 and 20 h. ( b ) Magnification of all XRD patterns and ( c ) fitting for the pTNTMZ10h in the 26–48 ◦ 2 θ range and around the (110) crystallographic planes for the both two bcc-1 and bcc-2 alloys. The corresponding crystallographic planes for each phase are marked between brackets. Dotted lines are used to highlight the shift in the XRD peaks. Metals 2020,10, 1463 7 of 15 When the milling time was increased, from 2 to 10 h, the intensity of Zr peaks gradually diminished and the corresponding peaks for the two bcc-1 and bcc-2 alloys underwent a shift to a lower 2theta degree. After 10 h of milling time, no Zr peaks were detected, as can be observed in Figure 1b. These inspections suggests that, during milling, Zr atoms (with the higher covalent radii) seemed to be constantly incorporated into the bcc matrices, and the two bcc alloys were completely formed after 10 h. Figure 1c shows the deconvolved XRD signal of the pTNTMZ10h pattern around the (110) bcc peaks, where the two bcc-1 and bcc-2 alloys were visible. At 20 h of milling time, the bcc-1 and bcc-2 alloys continued to shift slightly to a lower 2theta degree, suggesting a short-length redistribution of elements. However, at this longest milling time, the presence of Fe (coming from the stainless-steel milling media) was detected (bcc structure, Im3m SGS, COD no. 9008536). This presence is an undesirable aspect of the expected TiNbTaMoZr HEA as a potential biomaterial for a permanent bone-replacement implant. Iron is a biodegradable element [ 31 – 33 ], whose significant incorporation into the alloy might reduce mechanical performance due to the possible formation of porosity. This aspect constitutes a technological limitation that makes it unsuitable to apply longer milling times. Therefore, regarding the complete formation in two bcc-1 and bcc-2 alloys and the absence of the undesirable Fe, the pTNTMZ10h pattern was selected as the optimal synthesized specimen for its later consolidation. No oxidized metallic phases were indexed, which initially suggests that the mechanical alloying process was correctly carried out in the absence of oxygen. Subsequently, lattice parameters were calculated from the (110) major peaks of both bcc-1 and bcc-2 alloys. Their evolution throughout their milling time is shown in Figure 2. It can be confirmed how the lattice parameters increase as the alloying of the different elements progresses to the bcc structures. At 10 h of milling time, for pTNTMZ10h, chosen as the optimized as-synthesized specimen, the lattice parameters were calculated as 0.343 nm for the bcc-1 alloy and 0.319 nm for the bcc-2 alloy. Thus, the theoretical value, 0.335 nm, determined by the mixing law and by the different references in the COD of the bcc elemental transition metals (Ti (high-temperature allotrope, COD no. 9008554), a=3.3065 Å; Nb (COD no. 1539041), a =3.3063 Å; Ta (COD no. 1541266), a =3.3026 Å; Mo (COD no. 9008543), a =3.1473 Å) and Zr (high-temperature allotrope, COD no. 9008559), a =3.62 Å), lies between the corresponding values of the two bcc alloys. Metals 2020, 10, x FOR PEER REVIEW 7 of 16 When the milling time was increased, from 2 to 10 h, the intensity of Zr peaks gradually diminished and the corresponding peaks for the two bcc-1 and bcc-2 alloys underwent a shift to a lower 2theta degree. After 10 h of milling time, no Zr peaks were detected, as can be observed in Figure 1b. These inspections suggests that, during milling, Zr atoms (with the higher covalent radii) seemed to be constantly incorporated into the bcc matrices, and the two bcc alloys were completely formed after 10 h. Figure 1c shows the deconvolved XRD signal of the pTNTMZ10h pattern around the (110) bcc peaks, where the two bcc-1 and bcc-2 alloys were visible. At 20 h of milling time, the bcc-1 and bcc-2 alloys continued to shift slightly to a lower 2theta degree, suggesting a short-length redistribution of elements. However, at this longest milling time, the presence of Fe (coming from the stainless-steel milling media) was detected (bcc structure, Im3m SGS, COD no. 9008536). This presence is an undesirable aspect of the expected TiNbTaMoZr HEA as a potential biomaterial for a permanent bone-replacement implant. Iron is a biodegradable element [31–33], whose significant incorporation into the alloy might reduce mechanical performance due to the possible formation of porosity. This aspect constitutes a technological limitation that makes it unsuitable to apply longer milling times. Therefore, regarding the complete formation in two bcc-1 and bcc-2 alloys and the absence of the undesirable Fe, the pTNTMZ10h pattern was selected as the optimal synthesized specimen for its later consolidation. No oxidized metallic phases were indexed, which initially suggests that the mechanical alloying process was correctly carried out in the absence of oxygen. Subsequently, lattice parameters were calculated from the (110) major peaks of both bcc-1 and bcc-2 alloys. Their evolution throughout their milling time is shown in Figure 2. It can be confirmed how the lattice parameters increase as the alloying of the different elements progresses to the bcc structures. At 10 h of milling time, for pTNTMZ10h, chosen as the optimized as-synthesized specimen, the lattice parameters were calculated as 0.343 nm for the bcc-1 alloy and 0.319 nm for the bcc-2 alloy. Thus, the theoretical value, 0.335 nm, determined by the mixing law and by the different references in the COD of the bcc elemental transition metals (Ti (high-temperature allotrope, COD no. 9008554), a = 3.3065 Å; Nb (COD no. 1539041), a = 3.3063 Å; Ta (COD no. 1541266), a = 3.3026 Å; Mo (COD no. 9008543), a = 3.1473 Å) and Zr (high-temperature allotrope, COD no. 9008559), a = 3.62 Å), lies between the corresponding values of the two bcc alloys. 0 2 4 6 8 10 12 14 16 18 20 22 0.31 0.32 0.33 0.34 0.35 0.36 a th =0.335 nm Lattice parameter (nm) milling time (h) bcc-1 bcc-2 Figure 2. Experimental lattice parameters of bcc-1 and bcc-2 phases for all pTNTMZ specimens milled between 1 and 20 h. Dashed line shows the theoretical lattice parameter for a unique TiNbTaMoZr HEA, calculated by the mixing law. Figure 2. Experimental lattice parameters of bcc-1 and bcc-2 phases for all pTNTMZ specimens milled between 1 and 20 h. Dashed line shows the theoretical lattice parameter for a unique TiNbTaMoZr HEA, calculated by the mixing law. Metals 2020,10, 1463 8 of 15 RepresentativeSEMimagesareshownforpTNTMZ1h,pTNTMZ4h, pTNTMZ6handpTNTMZ10h (Figure 3), where a decrease in particle size with milling time is observed. Particles above 50 µ m were not visible when milling time reached 10 h. Indeed, the micrometric particle agglomerates that were detected were formed by submicrometric particles, as a typical morphology of powders synthesized by dry milling processes [24]. Metals 2020, 10, x FOR PEER REVIEW 8 of 16 Representative SEM images are shown for pTNTMZ1h, pTNTMZ4h, pTNTMZ6h and pTNTMZ10h (Figure 3), where a decrease in particle size with milling time is observed. Particles above 50 μm were not visible when milling time reached 10 h. Indeed, the micrometric particle agglomerates that were detected were formed by submicrometric particles, as a typical morphology of powders synthesized by dry milling processes [24]. Figure 3. SEM images for (a) pTNTMZ1h, (b) pTNTMZ4h, (c) pTNTMZ6h and (d) pTNTMZ10h showing the evolution of the morphology and size of particles with the milling time. To corroborate the element distribution of the various transition metals in the selected optimized as-synthesized pTNTMZ10h pattern, an SEM-EDS mapping was carried out (Figure 4). At first sight, good homogenization could be observed with the different transition metals distributed throughout the particles shown. Their calculated average atomic percentages (determined from 20 EDS points) indicate a nearly equiatomic alloy TiNb 0,91 Ta 0.92 Mo 0.90 Zr 0.97 (Ti = 21.3 ± 1.6; Nb = 19.4 ± 0.8; Ta = 19.5 ± 1.2; Mo = 19.1 ± 1.5; Zr = 20.7 ± 3.9). The amount of oxygen was practically negligible in all measurements, and always remained below the quantification limit of the SEM detector. Moreover, due to the presence of Ti, Nb, Ta, Mo and Ta throughout the particle, it is expected that these bcc alloys are, in fact, two TiNbTaMoZr-based HEAs. These results are in good agreement with previous work [26,27], where two bcc-TNTMZ alloys were also obtained. Figure 3. SEM images for ( a ) pTNTMZ1h, ( b ) pTNTMZ4h, ( c ) pTNTMZ6h and ( d ) pTNTMZ10h showing the evolution of the morphology and size of particles with the milling time. To corroborate the element distribution of the various transition metals in the selected optimized as-synthesized pTNTMZ10h pattern, an SEM-EDS mapping was carried out (Figure 4). At first sight, good homogenization could be observed with the different transition metals distributed throughout the particles shown. Their calculated average atomic percentages (determined from 20 EDS points) indicate a nearly equiatomic alloy TiNb 0,91 Ta 0.92 Mo 0.90 Zr 0.97 (Ti =21.3 ± 1.6; Nb =19.4 ± 0.8; Ta =19.5 ±1.2 ; Mo =19.1 ± 1.5; Zr =20.7 ± 3.9). The amount of oxygen was practically negligible in all measurements, and always remained below the quantification limit of the SEM detector. Moreover, due to the presence of Ti, Nb, Ta, Mo and Ta throughout the particle, it is expected that these bcc alloys are, in fact, two TiNbTaMoZr-based HEAs. These results are in good agreement with previous work [ 26 , 27 ], where two bcc-TNTMZ alloys were also obtained. Metals 2020,10, 1463 9 of 15 Metals 2020, 10, x FOR PEER REVIEW 9 of 16 Figure 4. SEM image and EDS-mapping of the pTNTMZ10h specimen showing the distribution of the different transition metals. (a) Secondary electron (SE) image, (b) titanium, (c) niobium, (d) tantalum, (e) molybdenum and (f) zirconium. An HRTEM study was carried out on the optimized pTNTMZ10h specimen (Figure 5) to corroborate the formation of the two bcc-1 and bcc-2 alloys detected by XRD. This study showed an HRTEM image for a particle and its corresponding FFT. In the annular electron diffraction pattern, two clear rings assigned to the (110) crystallographic planes for the bcc-1 and bcc-2 alloys can be observed, whereby the interplanar distances are 0.243 nm and 0.225 nm, respectively. Both crystallographic planes could also be observed in the lower HRTEM images of Figure 5, where the same interplanar distances could be calculated. Several other points were characterized in the FFT annular electron diffraction pattern, such as the (200) and (221) crystallographic planes for the majority bcc-2 alloy, whose interplanar distances were of 0.158 nm and 0.128 nm, respectively. These distances and crystallographic planes are in concordance with the most intense peaks found in the XRD patterns (Figure 1). These assertions, together with the absence of other diffraction spots, corroborate the unique formation of these two bcc-1 and bcc-2 TiNbTaZrMo alloys. Figure 4. SEM image and EDS-mapping of the pTNTMZ10h specimen showing the distribution of the different transition metals. ( a ) Secondary electron (SE) image, ( b ) titanium, ( c ) niobium, ( d ) tantalum, (e) molybdenum and (f) zirconium. AnHRTEMstudywascarriedoutontheoptimizedpTNTMZ10hspecimen(Figure5)tocorroborate the formation of the two bcc-1 and bcc-2 alloys detected by XRD. This study showed an HRTEM image for a particle and its corresponding FFT. In the annular electron diffraction pattern, two clear rings assigned to the (110) crystallographic planes for the bcc-1 and bcc-2 alloys can be observed, whereby the interplanar distances are 0.243 nm and 0.225 nm, respectively. Both crystallographic planes could also be observed in the lower HRTEM images of Figure 5, where the same interplanar distances could be calculated. Several other points were characterized in the FFT annular electron diffraction pattern, such as the (200) and (221) crystallographic planes for the majority bcc-2 alloy, whose interplanar distances were of 0.158 nm and 0.128 nm, respectively. These distances and crystallographic planes are in concordance with the most intense peaks found in the XRD patterns (Figure 1). These assertions, together with the absence of other diffraction spots, corroborate the unique formation of these two bcc-1 and bcc-2 TiNbTaZrMo alloys.