Microstructure and mechanical properties of sintered and heat-treated HfNbTaTiZr high entropy alloy

metals Article Microstructure and Mechanical Properties of Sintered and Heat-Treated HfNbTaTiZr High Entropy Alloy Jaroslav Málek 1,2,* , JiˇríZýka 1, František Lukáˇc 3,4 , Jakub ˇ Cížek 4, Lenka Kunˇcická5and Radim Kocich 6 1UJP PRAHA a.s., Nad Kamínkou 1345, 156 10 Prague-Zbraslav, Czech Republic; [email protected] 2CTU in Prague-Faculty of Mechanical Engineering, Karlovo Námˇestí13, 121 35 Praha 2, Czech Republic 3Institute of Plasma Physics CAS, Za Slovankou 3, 182 00 Praha 8, Czech Republic; [email protected] 4 Faculty of Mathematics and Physics, Charles University, V Holešoviˇck á ch 2, 180 00 Praha 8, Czech Republic; [email protected] 5Institute of Physics of Materials, CAS, Žižkova 22, 616 00 Brno, Czech Republic; [email protected] 6VŠB-Technical University of Ostrava, 17. Listopadu 15, 708 33 Ostrava 8, Czech Republic; [email protected] *Correspondence: [email protected]; Tel.:+420-7778-20414 Received: 25 October 2019; Accepted: 5 December 2019; Published: 7 December 2019   Abstract: High entropy alloys (HEAs) have attracted researchers’ interest in recent years. The aim of this work was to prepare the HfNbTaTiZr high entropy alloy via the powder metallurgy process and characterize its properties. The powder metallurgy process is a prospective solution for the synthesis of various alloys and has several advantages over arc melting (e.g., no dendritic structure, near net-shape, etc.). Cold isostatic pressing of blended elemental powders and subsequent sintering at 1400 ◦ C for various time periods up to 64 h was used. Certain residual porosity, as well as bcc2 (Nband Ta-rich) and hcp (Zrand Hf-rich) phases, remained in the bcc microstructure after sintering. The bcc2 phase was completely eliminated during annealing (1200 ◦ C/1h) and subsequent water quenching. The hardness values of the sintered specimens ranged from 300 to 400 HV10. The grain coarsening during sintering was significantly limited and the maximum average grain diameter after 64 h of sintering was approximately 60 µ m. The compression strength at 800 ◦ C was 370 MPa and decreased to 47 MPa at 1200 ◦ C. Porosity can be removed during the hot deformation process, leading to an increase in hardness to ~450 HV10. Keywords: high entropy alloys; sintering; microstructure 1. Introduction The concept of the so-called “high entropy alloys” (HEAs), also known as “multi-principal element alloys” (MPEAs) or “complex concentrated alloys” (CCAs)[ 1 ], has been attracting attention since its formulation in the early 2000s by Yeh et al. [ 2 ] and Cantor et al. [ 3 ]. The concept of HEAs is different from other alloys of modern structural materials (e.g., composites) which may also consist of multiple metallic elements [ 4 , 5 ]. The traditional design of new alloys with one principal element is replaced by a new group of materials with multiple elements. HEAs typically contain at least five elements with equimolar or near equimolar (5–35 at.%) concentrations [ 1 , 6 ]. Numerous parameters were proposed by various authors in order to predict whether an alloy will exhibit a microstructure consisting of a solid solution with no intermetallic phases [ 1 , 7 – 10 ]. The matrixes of high entropy alloys are typically bcc or fcc, but recently even HEAs with the hcp matrix have been reported [ 1 , 11 – 14 ]. The HEAs with the bcc matrix usually exhibit high strength and lower plasticity whereas fcc HEAs have lower strength and high plasticity [1,6]. Metals 2019,9, 1324; doi:10.3390/met9121324 www.mdpi.com/journal/metals Metals 2019,9, 1324 2 of 18 HEAs exhibit intriguing combinations of properties and structural phenomena; high lattice distortions or sluggish diffusion effects are frequently reported (among others [ 6 ]). Such properties are usually attributed to their high entropy and its related effects. As a result of these phenomena, HEAs have the potential to be used in many applications (e.g., high temperature materials, cryogenic materials, etc.) [ 1 , 6 ]. The HfNbTaTiZr equimolar alloy belongs to a bcc structure family of HEAs. It is a promising high entropy alloy with advantageous prospective usage. Given its composition consisting of refractory metals with melting points above 1850 ◦ C [ 1 , 15 , 16 ], this particular HEA is a promising refractory alloy which can possibly be used to substitute tungsten-based alloys in selected applications [ 17 ].This alloy exhibits quite high tensile strength at room temperature. However, the values reported by various authors are broadly scattered. Dirras et al. [ 18 ] reported a tensile strength of 850 MPa, as well as Schuh et al. [ 19 ], reaching even 1900 MPa after cold deformation. Those properties were strongly influenced by the microstructure. Senkov et al. [ 20 ] showed that the tensile strength of this alloy may reach 1150 MPa by retaining ~15% of elongation after suitable thermo-mechanical processing. Very good combinations of high tensile strength and elongation have also been reported by other authors [ 21 – 23 ]. Excellent compression strength values at elevated temperatures ( i.e., 535 MPa at 800 ◦ C or 295 MPa at 1000 ◦ C) were also reported by Senkov et. al. [ 16 ].The HfNbTaTiZr alloy contains only biocompatible elements. Its biocompatibility presupposes the alloy also for biomedical applications [ 24 – 26 ]. Therefore, HfNbTaTiZr can be considered as a modern biocompatible material, together with other innovative titanium-based alloys [27,28]. HEAs are usually prepared via arc melting processes, which often results in an inhomogeneous dendritic microstructure [ 15 , 18 , 22 , 29 – 40 ]. The dendritic microstructures are usually removed by subsequent heat treatment. However, in some cases, the dendritic structure is not fully removed after annealing and significant grain coarsening can also be introduced during annealing [ 16 , 31 , 33 , 38 ]. It has been reported that large grain size is not desirable for most (not for all) applications as, in most cases, fine grains improve strength and ductility [ 22 , 23 ]. Nevertheless, modern manufacturing methods, such as powder metallurgy technologies [ 41 ] and specialized deformation methods (methods of severe plastic deformation—SPD) [ 42 ], can be used to refine the structures and enhance the mechanical properties of the processed materials. Powder metallurgy is a prospective method for the production of metallic materials as it may overcome some of the negative aspects of arc melting (e.g., coarse grains, dendritic structure) [ 43 ]. Spark plasma sintering (SPS) and metal injection moulding (MIM) are the most frequently used powder metallurgy technologies, along with others. Only a few works using the powder metallurgy process for the synthesis of HfNbTATiZr (or similar) HEAs have been reported to date (e.g., [ 9 , 43 ]). Our previous works reported on the preparation of the HfNbTaTiZr equimolar alloy via the arc melting [ 44 – 46 ] and SPS [ 47 ] processes. However, the disadvantage of the SPS technology is complicated powder preparation. The powder of the HfNbTaTiZr alloy is currently not commercially available and it is difficult (and expensive) to obtain—usually from gas atomization of the arc melted alloy. Due to these reasons, the present work is focused on the preparation of the HfNbTATiZr alloy from elemental powders which are commercially available. The aim of this work was to synthesize the HfNbTaTiZr alloy via the blended elemental powder metallurgy process and to evaluate the influence of the process parameters (mainly sintering time) on the microstructure and mechanical properties of the final product. 2. Materials and Methods The samples of the HfNbTaTiZr alloy were prepared via the powder metallurgy process. Elemental powders (granularity: 325 mesh, i.e., grain size <44 µ m; purity: min. 99.5%) with irregular shape (HDH powders) were supplied by Huarui Industrial Materials Co., Ltd. (Chengdu, China). The powders were weighted to obtain the desired chemical composition of the equimolar HfNbTaTiZr alloy, mixed in a Turbula 2F device(WAB GmbH, Basel, Switzerland), and filled into the cylinder-shaped molds (inner diameter: 20 mm). These preparation steps were performed under an Ar protective Metals 2019,9, 1324 3 of 18 atmosphere. The specimens were cold isostatically pressed (CIP) using a pressure of 400 MPa and subsequently sintered. The sintering was performed in a vacuum (better than 10 −3 Pa) furnace at 1400 ◦ C (with 20 ◦ C/min heating rate) for 2, 4, 8, 16, 32 and 64 h (corresponding samples were denoted as “2 h”, “4 h”, “8 h”, “16h”, “32h” and “64h”, respectively). The sintered specimens were further annealed under the argon protective atmosphere at 1200 ◦ C for 1 h and subsequently water quenched. The latter specimens were denoted by the sintering period and a suffix “1200 ◦C/1h/WQ”. The microstructure of the specimens was studied by light microscopy (LM) using a Nikon EPIPHOT 3000 microscope (Nikon, Melville, New York, NY, USA) and scanning electron microscopy (SEM) using a JEOL 7650F (JEOL, Akishima, Tokyo, Japan)scanning electron microscope equipped with an electron back scatter diffraction (EBSD) detector Nordlys II (Oxford Instruments, Abingdon on Thames, UK). The SEM observations were done in back-scattered electron mode (BSE) unless otherwise noted. NIS Elements software (Nikon, Melville, New York, NY, USA) was used for image analysis. The specimens for LM and SEM observations were prepared by a standard metallographic process (ground up to #4000 with SiC papers and polished with Struers OP-S emulsion with the addition of H 2 O 2 ). For etching, 3 mL of HF +8 mL of HNO 3 +100 mL of H 2 O etchants were used. Some specimens were also observed using a JEOL 2000 EX (JEOL, Akishima, Tokyo, Japan) transmission electron microscope (TEM) to provide a more detailed microstructure characterization. The specimens for TEM were ground to a thickness of ~100 µ m and appropriate dimensions for the JEOL EM-09100 Ion Slicer (JEOL, Akishima, Tokyo, Japan), where the specimens were thinned to a final thickness. The phase identification was also carried out using X-ray diffraction analysis (XRD) on a Bruker D8 Discover diffractometer (Bruker, Karlsruhe, Germany) using a Cu anode divergent beam aimed at the polished cross section of a bar. Quantitative Rietveld refinement analysis was performed by TOPAS V5 software (Bruker, Karlsruhe, Germany). The Differential Scanning Calorimetry (DSC) method was carried out on Setaram Labsys Evo apparatus (Setaram, Caluire, France) in order to obtain information about the phase stability. The samples were heated in Al 2 O 3 crucibles with a heating rate of 7 ◦ C/min in an Ar flow of 100 mL/min. The Vickers hardness was determined using a Zwick/Roell ZHU 250 top hardness tester (Zwick/Roell, Ulm, Germany) with a load of 98.1 N (according to the ISO 6507 standard). At least seven values were determined for each measurement. The compression tests were carried out using Zwick furnaces (Zwick/Roell, Ulm, Germany) in the Ar protective atmosphere at the temperatures of 800, 900, 1000, 1100 and 1200 ◦ C and a strain rate of 10 −4 . The samples for the compression test were 10 mm long and had a circular cross-section with the area of 80 mm 2 . Given the relatively low strain rate, the samples had the tendency to exhibit creep behavior. Due to this, the compression tests were all terminated after 90 min of load exposure. The oxygen content in the sintered specimens was determined by a Bruker Galileo G8 gas fusion analyzer (Bruker, Karlsruhe, Germany). At least three measurements for each specimen were performed. The grain size was measured using the line intercept method from at least five images from different regions of the specimen. The images covered a minimum area of 800 µ m × 600 µ m. The image analysis was used to determine the porosity (the same images as for grain size determination were used). 3. Results The microstructure of the processed specimens is described first, followed by the presentation of the mechanical properties which are supposed to correlate with the microstructure. 3.1. Microstructure The specimens exhibited certain residual porosity after sintering (determined by the image analysis of at least five light micrographs for each value).The results are shown in Figure 1. The pores had typical dimensions up to 50 µm. Most of the pores were between 10 and 30 µm. Metals 2019,9, 1324 4 of 18 Metals 2019, 9, x FOR PEER REVIEW 4 of 17 Figure 1.The development of porosity vs. sintering time at 1400°C. The oxygen content was measured for the sintered and as-pressed (sintering time 0h) specimens. A slightly increasing oxygen content with an increasing sintering time can be seen in Figure 2 (despite the large experimental scatter). Figure 2.Measured oxygen content vs. sintering time in as-sintered specimens. The microstructure of the as-sintered specimens underwent changes during the sintering process. The original mixture of pressed elemental powders (Figure 3a) was dissolved and formed HfNbTaTiZr alloy (Figure 3b–d). The matrix had a chemical composition close to the equimolar composition in all sintered specimens (see Table 1). Numerous chemical heterogeneities in the “2h”, “4h” and “8h” specimens (highlighted in the “2h” specimen—Figure 3b) could be clearly observed. Their occurrence diminished with an increasing sintering time and in the specimens sintered for 16 h (or more) could only be rarely seen (Figures3c,d). The heterogeneities were identified as areas with an increased Ta content. However, their number and volume fraction were lower than the fraction of Ta powder in the original powder blend. No areas with significantly higher Nb (or other elements except forTa) content that could be ascribed to insufficient dissolution of the original powder particles were observed in any of the sintered specimens. According to EBSD analysis (not shown here), neither the sintered nor annealed specimens exhibited any preferred grain orientation (texture). Table 1.Average chemical composition (in at.%) of observed phases indicated in Figure 3e. Phase Ti Zr Nb Hf Ta Matrix bcc 19.9 ± 0.5 19.0 ± 1.3 21.1 ± 1.6 20.2 ± 0.9 19.8 ± 1.7 precipitate bcc2 22.0 ± 2.1 7.0 ± 1.4 28.3 ± 2.8 11.1 ± 1.2 31.3 ± 3.4 precipitate hcp 7.7 ± 1.5 40.7 ± 1.6 6.5 ± 2.5 38.8 ± 3.0 7.8 ± 3.0 Figure 1. The development of porosity vs. sintering time at 1400 ◦C. The oxygen content was measured for the sintered and as-pressed (sintering time 0h) specimens. AslightlyincreasingoxygencontentwithanincreasingsinteringtimecanbeseeninFigure2(despitethe large experimental scatter). Metals 2019, 9, x FOR PEER REVIEW 4 of 17 Figure 1.The development of porosity vs. sintering time at 1400°C. The oxygen content was measured for the sintered and as-pressed (sintering time 0h) specimens. A slightly increasing oxygen content with an increasing sintering time can be seen in Figure 2 (despite the large experimental scatter). Figure 2.Measured oxygen content vs. sintering time in as-sintered specimens. The microstructure of the as-sintered specimens underwent changes during the sintering process. The original mixture of pressed elemental powders (Figure 3a) was dissolved and formed HfNbTaTiZr alloy (Figure 3b–d). The matrix had a chemical composition close to the equimolar composition in all sintered specimens (see Table 1). Numerous chemical heterogeneities in the “2h”, “4h” and “8h” specimens (highlighted in the “2h” specimen—Figure 3b) could be clearly observed. Their occurrence diminished with an increasing sintering time and in the specimens sintered for 16 h (or more) could only be rarely seen (Figures3c,d). The heterogeneities were identified as areas with an increased Ta content. However, their number and volume fraction were lower than the fraction of Ta powder in the original powder blend. No areas with significantly higher Nb (or other elements except forTa) content that could be ascribed to insufficient dissolution of the original powder particles were observed in any of the sintered specimens. According to EBSD analysis (not shown here), neither the sintered nor annealed specimens exhibited any preferred grain orientation (texture). Table 1.Average chemical composition (in at.%) of observed phases indicated in Figure 3e. Phase Ti Zr Nb Hf Ta Matrix bcc 19.9 ± 0.5 19.0 ± 1.3 21.1 ± 1.6 20.2 ± 0.9 19.8 ± 1.7 precipitate bcc2 22.0 ± 2.1 7.0 ± 1.4 28.3 ± 2.8 11.1 ± 1.2 31.3 ± 3.4 precipitate hcp 7.7 ± 1.5 40.7 ± 1.6 6.5 ± 2.5 38.8 ± 3.0 7.8 ± 3.0 Figure 2. Measured oxygen content vs. sintering time in as-sintered specimens. The microstructure of the as-sintered specimens underwent changes during the sintering process. The original mixture of pressed elemental powders (Figure 3a) was dissolved and formed HfNbTaTiZr alloy (Figure 3b–d). The matrix had a chemical composition close to the equimolar composition in all sintered specimens (see Table 1). Numerous chemical heterogeneities in the “2h”, “4h” and “8h” specimens (highlighted in the “2h” specimen—Figure 3b) could be clearly observed. Their occurrence diminished with an increasing sintering time and in the specimens sintered for 16 h (or more) could only be rarely seen (Figure 3c,d). The heterogeneities were identified as areas with an increased Ta content. However, their number and volume fraction were lower than the fraction of Ta powder in the original powder blend. No areas with significantly higher Nb (or other elements except forTa) content that could be ascribed to insufficient dissolution of the original powder particles were observed in any of the sintered specimens. According to EBSD analysis (not shown here), neither the sintered nor annealed specimens exhibited any preferred grain orientation (texture). Table 1. Average chemical composition (in at.%) of observed phases indicated in Figure 3e. Phase Ti Zr Nb Hf Ta Matrix bcc 19.9 ±0.5 19.0 ±1.3 21.1 ±1.6 20.2 ±0.9 19.8 ±1.7 precipitate bcc2 22.0 ±2.1 7.0 ±1.4 28.3 ±2.8 11.1 ±1.2 31.3 ±3.4 precipitate hcp 7.7 ±1.5 40.7 ±1.6 6.5 ±2.5 38.8 ±3.0 7.8 ±3.0 Metals 2019,9, 1324 5 of 18 Metals 2019, 9, x FOR PEER REVIEW 5 of 17 (a) (b) (c) (d) (e) Figure 3.Microstructure of specimens shown in scanning electron microscopy (SEM)—back-scattered electron (BSE) mode: (a) pressed powders before sintering; (b) sintered 1400°C/2h—areas with increased Ta content are marked by arrows; (c) sintered 1400°C/16h; (d) sintered 1400°C/64h—coarse hcp phase marked by arrow; (e) sintering 1400°C/8h. Numerous precipitates can be seen in the microstructure of the sintered specimens (Figures 3–5). An example of the XRD pattern of sample “64h” is plotted in Figure 6a. The XRD analysis confirmed the presence of two bcc and one hcp phases. Additionally, a very small fraction (<1 wt.%) of (Hf, Zr)O 2 oxides was detected in the “64h” specimen. This could be ascribed either to the local contamination of the specimen or to a higher oxygen content (see Figure 2) which may lead to (Hf, Zr)O 2 formation. The dominating phase, i.e., matrix, denoted here as “bcc”, has the bcc structure with the lattice parameter a = 3.399(8) Å and its chemical composition is close to the equimolar concentration HfNbTaTiZr (see Table 1). The bcc matrix contains precipitates which have either a bcc or a hcp structure. The precipitates with a bcc structure, denoted “bcc2”, exhibit a lattice parameter a = 3.347(3) Å, which is lower than that of the bcc matrix. According to EDS analysis, the bcc2 precipitates are enriched in Ta and Nb (see Table 1). The precipitates with hcp structure, denoted “hcp”, are characterized by lattice parameters a = 3.205(3) Å and c = 5.113(7) Å and are rich Figure 3. Microstructure of specimens shown in scanning electron microscopy (SEM)—back-scattered electron (BSE) mode: ( a ) pressed powders before sintering; ( b ) sintered 1400 ◦ C/2h—areas with increased Ta content are marked by arrows; ( c ) sintered 1400 ◦ C/16h; ( d ) sintered 1400 ◦ C/64h—coarse hcp phase marked by arrow; (e) sintering 1400 ◦C/8h. Numerous precipitates can be seen in the microstructure of the sintered specimens (Figures 3–5). An example of the XRD pattern of sample “64h” is plotted in Figure 6a. The XRD analysis confirmed the presence of two bcc and one hcp phases. Additionally, a very small fraction (<1 wt.%) of (Hf, Zr)O 2 oxides was detected in the “64h” specimen. This could be ascribed either to the local contamination of the specimen or to a higher oxygen content (see Figure 2) which may lead to (Hf, Zr)O 2 formation. The dominating phase, i.e., matrix, denoted here as “bcc”, has the bcc structure with the lattice parameter a=3.399(8) Å and its chemical composition is close to the equimolar concentration HfNbTaTiZr (see Table 1). The bcc matrix contains precipitates which have either a bcc or a hcp structure. The precipitates with a bcc structure, denoted “bcc2”, exhibit a lattice parameter a=3.347(3) Å, which is lower than that of the bcc matrix. According to EDS analysis, the bcc2 precipitates are enriched in Ta and Nb (see Table 1). The precipitates with hcp structure, denoted “hcp”, are characterized by Metals 2019,9, 1324 6 of 18 lattice parameters a=3.205(3) Å and c=5.113(7) Å and are rich in Hf and Zr (see Table 1). It was determined by XRD analysis that the fraction of the Zrand Hf-rich phase (hcp) is lower than that of the Taand Nb-rich phase (bcc2) in the sintered specimens (see Table 2), which is consistent with the SEM observations (Figure 3d,e). One type of precipitate is located both in the vicinity of the grain boundaries and in the grain interiors (see Figure 3d,e). The EDS analysis revealed that both of these precipitates are Taand Nb-rich phases (i.e., bcc2 phase; see Table 1). Other precipitates of irregular shapes could be observed along the grain boundaries (also depicted with arrows in Figure 3e). These precipitates were typically smaller in size (up to 1 µ m). They were indexed as hcp phase according to TEM-SAED analysis (Figure 5) and were enriched with Zr and Hf (see Table 1). It has to be mentioned that, in some cases (especially for the hcp phase), the precipitates were so fine that the EDS signal could partly originate from the surrounding area. Due to this reason, more measurements were taken throughout all the specimens and the values summarized in Table 1are the average chemical compositions. Additionally, the significantly coarser (few tens of µ m) hcp precipitates (Hfand Zr-rich) were observed along the edge of the specimens (Figure 4) and locally also inside the “64h” specimen (see Figure 3d depicted by a white arrow). These precipitates were of a similar irregular shape to the fine ones. According to TEM-SAED analysis, the hcp precipitates had an orientation relationship with the matrix {110}matrix//{0001}hcp and [001]matrix//h2110ihcp (Figure 5). Table 2. Weight fractions of phases detected by XRD. Specimen bcc hcp (Zrand Hf-rich) bcc2 (Taand Nb-rich) “2 h” 68.5% 5% 26.5% “4 h” 59.5% 11.3% 29.2% “8 h” 65.6% 11.6% 22.8% “16 h” 65.0% 6.8% 28.2% “32 h” 58.5% 13% 28.5% “64 h” * 52.2% 5.3% 41.3% “2 h” +1200 ◦C/1 h/WQ 87.5% 12.5% - “4 h” +1200 ◦C/1 h/WQ 91.8% 8.2% - “8 h” +1200 ◦C/1 h/WQ 95.3% 4.7% - “16 h” +1200 ◦ C/1 h/WQ 94.4% 5.6% - “32 h” +1200 ◦ C/1 h/WQ 94.6% 5.4% - “64 h” +1200 ◦ C/1 h/WQ 92.1% 7.9% - * This specimen also contained a very small fraction of (Hf, Zr)O2oxide. Metals 2019, 9, x FOR PEER REVIEW 6 of 17 in Hf and Zr (see Table 1). It was determined by XRD analysis that the fraction of the Zrand Hf-rich phase (hcp) is lower than that of the Taand Nb-rich phase (bcc2) in the sintered specimens (see Table 2), which is consistent with the SEM observations (Figure 3d,e). One type of precipitate is located both in the vicinity of the grain boundaries and in the grain interiors (see Figure 3d,e). The EDS analysis revealed that both of these precipitates are Taand Nb-rich phases (i.e., bcc2 phase; see Table1). Other precipitates of irregular shapes could be observed along the grain boundaries (also depicted with arrows in Figure 3e). These precipitates were typically smaller in size (up to 1 μm). They were indexed as hcp phase according to TEM-SAED analysis (Figure 5) and were enriched with Zr and Hf (see Table 1). It has to be mentioned that, in some cases (especially for the hcp phase), the precipitates were so fine that the EDS signal could partly originate from the surrounding area. Due to this reason, more measurements were taken throughout all the specimens and the values summarized in Table1 are the average chemical compositions. Additionally, the significantly coarser (few tens of μm) hcp precipitates (Hfand Zr-rich) were observed along the edge of the specimens (Figure 4) and locally also inside the “64h” specimen (see Figure 3d depicted by a white arrow). These precipitates were of a similar irregular shape to the fine ones. According to TEM-SAED analysis, the hcp precipitates had an orientation relationship with the matrix {110}matrix//{0001}hcp and [001]matrix//󰇟21 1 0󰇠hcp (Figure 5). Table 2.Weight fractions of phases detected by XRD. Specimen bcc hcp (Zrand Hf-rich) bcc2 (Taand Nb-rich) “2 h” 68.5% 5% 26.5% “4 h” 59.5% 11.3% 29.2% “8 h” 65.6% 11.6% 22.8% “16 h” 65.0% 6.8% 28.2% “32 h” 58.5% 13% 28.5% “64 h” * 52.2% 5.3% 41.3% “2 h” + 1200 °C/1 h/WQ 87.5% 12.5% - “4 h” + 1200 °C/1 h/WQ 91.8% 8.2% - “8 h” + 1200 °C/1 h/WQ 95.3% 4.7% - “16 h” + 1200 °C/1 h/WQ 94.4% 5.6% - “32 h” + 1200 °C/1 h/WQ 94.6% 5.4% - “64 h” + 1200 °C/1 h/WQ 92.1% 7.9% - * This specimen also contained a very small fraction of (Hf, Zr)O2 oxide. Figure 4.Microstructure after sintering at 1400°C/64h. The specimen surface area at the bottom (LM). Figure 4. Microstructure after sintering at 1400 ◦ C/64h. The specimen surface area at the bottom (LM). Metals 2019,9, 1324 7 of 18 Metals 2019, 9, x FOR PEER REVIEW 7 of 17 Figure 5.Microstructure with hcp phase after sintering at 1400°C/16h with the diffraction pattern (TEM-SAED). Figure 6.X-ray diffraction (XRD) pattern for (a) sample “64h”; (b) sample “64h+1200°C/1h/WQ“. An example of the XRD pattern of sample “64h+1200°C/1h/WQ” is plotted in Figure 6b. In contrast to the as-sintered samples, no bcc2 phase was detected in the samples annealed at 1200°C for 1 h and water quenched. The XRD analysis of the 1200°C/1h/WQ samples revealed only the presence of a single bcc phase (matrix) with slightly broadened peaks and a certain fraction of the hcp (Hf, Zr-rich) phase (Table2).The microstructure of the quenched samples consisted of equiaxed grains with numerous fine precipitates located both along the grain boundaries and inside the grain interiors (see Figure 7). The TEM analysis (see Figure 8) proved that the fine precipitates in the solution-treated specimens were hcp particles with the same orientation relationship {110}matrix//{0001}hcp and [001]matrix//󰇟21 1 0󰇠hcp as in the as-sintered specimens. Water quenching also resulted in the emergence of cracks within some specimens, which indicated a high internal stress caused by the high cooling rate and limited plasticity at room temperature. Figure 5. Microstructure with hcp phase after sintering at 1400 ◦ C/16h with the diffraction pattern (TEM-SAED). Metals 2019, 9, x FOR PEER REVIEW 7 of 17 Figure 5.Microstructure with hcp phase after sintering at 1400°C/16h with the diffraction pattern (TEM-SAED). Figure 6.X-ray diffraction (XRD) pattern for (a) sample “64h”; (b) sample “64h+1200°C/1h/WQ“. An example of the XRD pattern of sample “64h+1200°C/1h/WQ” is plotted in Figure 6b. In contrast to the as-sintered samples, no bcc2 phase was detected in the samples annealed at 1200°C for 1 h and water quenched. The XRD analysis of the 1200°C/1h/WQ samples revealed only the presence of a single bcc phase (matrix) with slightly broadened peaks and a certain fraction of the hcp (Hf, Zr-rich) phase (Table2).The microstructure of the quenched samples consisted of equiaxed grains with numerous fine precipitates located both along the grain boundaries and inside the grain interiors (see Figure 7). The TEM analysis (see Figure 8) proved that the fine precipitates in the solution-treated specimens were hcp particles with the same orientation relationship {110}matrix//{0001}hcp and [001]matrix//󰇟21 1 0󰇠hcp as in the as-sintered specimens. Water quenching also resulted in the emergence of cracks within some specimens, which indicated a high internal stress caused by the high cooling rate and limited plasticity at room temperature. Figure 6. X-ray diffraction (XRD) pattern for (a) sample “64h”; (b) sample “64h+1200 ◦C/1h/WQ”. An example of the XRD pattern of sample “64h+1200 ◦ C/1h/WQ” is plotted in Figure 6b. In contrast to the as-sintered samples, no bcc2 phase was detected in the samples annealed at 1200 ◦ C for 1 h and water quenched. The XRD analysis of the 1200 ◦ C/1h/WQ samples revealed only the presence of a single bcc phase (matrix) with slightly broadened peaks and a certain fraction of the hcp (Hf, Zr-rich) phase (Table 2).The microstructure of the quenched samples consisted of equiaxed grains with numerous fine precipitates located both along the grain boundaries and inside the grain interiors (see Figure 7). The TEM analysis (see Figure 8) proved that the fine precipitates in the solution-treated specimens were hcp particles with the same orientation relationship {110} matrix //{0001} hcp and [001] matrix // h2110ihcp as in the as-sintered specimens. Water quenching also resulted in the emergence of cracks within some specimens, which indicated a high internal stress caused by the high cooling rate and limited plasticity at room temperature. Metals 2019,9, 1324 8 of 18 Metals 2019, 9, x FOR PEER REVIEW 8 of 17 Figure 7.Microstructure of the 64h + 1200°C/1h/WQ sample with a coarse and fine hcp phase (LM). Figure 8.Microstructure of the 64h + 1200°C/1h/WQ sample with a hcp phase (TEM-SAED). The average grain size is plotted in Figure 9 as a function of the sintering time at 1400°C. The mean grain size increased with the sintering period from 25 μm after 2 h of sintering to 60 μm after 64 h of sintering. The grain size of each specimen remained unchanged (with respect to the experimental scatter) during 1200°C/1h annealing. Figure 9.Grain size vs. sintering time. Experimental data are plotted by circles; the development of grain size calculated using Equation (1) (see in chapter 4. Discussion) for Q = 363 kJ/mol and Q = 480 kJ/mol is plotted by solid lines. The DSC curves for the sintered samples heated at the heating rate of 7°C/min (Figure 10) show a significant endothermic peak starting at approximately 1050°C with a maximum at around 1300–1350°C. The endothermic peak can be attributed to the dissolution of the secondary phases (bcc2 and hcp). In order to find the origin of the endothermic peak, DSC measurements were also taken of the samples annealed at 1200°C and quenched in water. The DSC curve of the Figure 7. Microstructure of the 64h +1200 ◦C/1h/WQ sample with a coarse and fine hcp phase (LM). Metals 2019, 9, x FOR PEER REVIEW 8 of 17 Figure 7.Microstructure of the 64h + 1200°C/1h/WQ sample with a coarse and fine hcp phase (LM). Figure 8.Microstructure of the 64h + 1200°C/1h/WQ sample with a hcp phase (TEM-SAED). The average grain size is plotted in Figure 9 as a function of the sintering time at 1400°C. The mean grain size increased with the sintering period from 25 μm after 2 h of sintering to 60 μm after 64 h of sintering. The grain size of each specimen remained unchanged (with respect to the experimental scatter) during 1200°C/1h annealing. Figure 9.Grain size vs. sintering time. Experimental data are plotted by circles; the development of grain size calculated using Equation (1) (see in chapter 4. Discussion) for Q = 363 kJ/mol and Q = 480 kJ/mol is plotted by solid lines. The DSC curves for the sintered samples heated at the heating rate of 7°C/min (Figure 10) show a significant endothermic peak starting at approximately 1050°C with a maximum at around 1300–1350°C. The endothermic peak can be attributed to the dissolution of the secondary phases (bcc2 and hcp). In order to find the origin of the endothermic peak, DSC measurements were also taken of the samples annealed at 1200°C and quenched in water. The DSC curve of the Figure 8. Microstructure of the 64h +1200 ◦C/1h/WQ sample with a hcp phase (TEM-SAED). The average grain size is plotted in Figure 9as a function of the sintering time at 1400 ◦ C. The mean grain size increased with the sintering period from 25 µ m after 2 h of sintering to 60 µ m after 64 h of sintering. The grain size of each specimen remained unchanged (with respect to the experimental scatter) during 1200 ◦C/1h annealing. Metals 2019, 9, x FOR PEER REVIEW 8 of 17 Figure 7.Microstructure of the 64h + 1200°C/1h/WQ sample with a coarse and fine hcp phase (LM). Figure 8.Microstructure of the 64h + 1200°C/1h/WQ sample with a hcp phase (TEM-SAED). The average grain size is plotted in Figure 9 as a function of the sintering time at 1400°C. The mean grain size increased with the sintering period from 25 μm after 2 h of sintering to 60 μm after 64 h of sintering. The grain size of each specimen remained unchanged (with respect to the experimental scatter) during 1200°C/1h annealing. Figure 9.Grain size vs. sintering time. Experimental data are plotted by circles; the development of grain size calculated using Equation (1) (see in chapter 4. Discussion) for Q = 363 kJ/mol and Q = 480 kJ/mol is plotted by solid lines. The DSC curves for the sintered samples heated at the heating rate of 7°C/min (Figure 10) show a significant endothermic peak starting at approximately 1050°C with a maximum at around 1300–1350°C. The endothermic peak can be attributed to the dissolution of the secondary phases (bcc2 and hcp). In order to find the origin of the endothermic peak, DSC measurements were also taken of the samples annealed at 1200°C and quenched in water. The DSC curve of the Figure 9. Grain size vs. sintering time. Experimental data are plotted by circles; the development of grain size calculated using Equation (1) (see in chapter 4. Discussion) for Q=363 kJ/mol and Q=480 kJ/mol is plotted by solid lines. The DSC curves for the sintered samples heated at the heating rate of 7 ◦ C/min (Figure 10) show a significant endothermic peak starting at approximately 1050 ◦ C with a maximum at around 1300–1350 ◦ C. The endothermic peak can be attributed to the dissolution of the secondary phases ( bcc2 and hcp ). In order to find the origin of the endothermic peak, DSC measurements were also taken Metals 2019,9, 1324 9 of 18 of the samples annealed at 1200 ◦ C and quenched in water. The DSC curve of the “16h+1200 ◦ C/1h/WQ” sample plotted in Figure 10c exhibits a similar endothermic peak to the sintered alloys (Figure 10a,b). Note that the DSC curve of the sample “64h” contains two separated endothermic peaks. The peak in the range 1250–1450 ◦ C, ascribed to dissolution of the hcp phase, is preceded by an endothermic peak occurring in the temperature range 1050–1250 ◦C (see Figure 10d). Metals 2019, 9, x FOR PEER REVIEW 9 of 17 “16h+1200°C/1h/WQ” sample plotted in Figure 10c exhibits a similar endothermic peak to the sintered alloys (Figure 10a,b). Note that the DSC curve of the sample “64h” contains two separated endothermic peaks. The peak in the range 1250–1450°C, ascribed to dissolution of the hcp phase, is preceded by an endothermic peak occurring in the temperature range 1050–1250°C (see Figure 10d). Figure 10.DSC curves for the samples heated at the heating rate of 7°C/min. The vertical dashed line indicates the onset of the endothermic peak ascribed to the dissolution of the hcp phase. (a) 4 h, (b) 8 h, (c) 16 h + 1200 °C/1 h/WQ, (d) 64 h. 3.2. Mechanical Properties The hardness increases with the increasing sintering time period, as can be seen in Figure 11. The hardness values have a relatively large experimental scatter, which is probably caused by the residual porosity, and partly also by chemical inhomogeneities (especially within specimens after shorter sintering periods). Figure 10. DSC curves for the samples heated at the heating rate of 7 ◦ C/min. The vertical dashed line indicates the onset of the endothermic peak ascribed to the dissolution of the hcp phase. ( a ) 4 h, ( b ) 8 h, (c) 16 h +1200 ◦C/1 h/WQ, (d) 64 h. 3.2. Mechanical Properties The hardness increases with the increasing sintering time period, as can be seen in Figure 11. The hardness values have a relatively large experimental scatter, which is probably caused by the residual porosity, and partly also by chemical inhomogeneities (especially within specimens after shorter sintering periods). Metals 2019,9, 1324 16 of 18 20. Senkov, O.N.; Pilchak, A.L.; Semiatin, S.L. 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