*3.2. Spark Plasma Sintering*

Contamination of samples sintered at 750 ◦C and 850 ◦C is shown in Table 2. The content of oxygen in the initial powder slightly increased after sintering, while in the milled powder, it decreased due to the decomposition of stearic acid at high temperatures. The sintering temperature did not have a significant effect on contamination by oxygen for both initial and milled powders.


**Table 2.** Contamination of sintered samples by oxygen.

Figure 4 shows BSE micrographs of sintered powders. Black areas in these BSE figures are α-phase, pores and impurities from polishing. In order to analyze images, these three phenomena were distinguished using SE observation. SE micrographs of the same areas as in Figure 4 are shown in Figure 5. In SE signal, pores are black, impurities from polishing are white and the α-phase is grey.

SEM micrographs of initial powder sintered at 750 ◦C and at 850 ◦C are shown in Figures 4 and 5. Initial powder sintered at 750 ◦C (Figures 4a and 5a) contains a significant amount of α-phase while initial powder sintered at 850 ◦C (cf. Figures 4b and 5b) contains only a small amount of grain boundary α-phase and black areas in BSE micrographs are mostly impurities from polishing. BSE micrographs of milled powder sintered at 750 ◦C and at 850 ◦C are shown in Figure 4c,d, respectively and in Figure 5c,d, respectively. Sintered milled powder has finer microstructure and contains higher amount of α-phase in comparison with the initial one. With increasing sintering temperature the microstructure coarsens and the fraction of α-phase decreases.

**Figure 4.** Examples of back-scattered electrons (BSE) micrographs (**a**) initial powder sintered at 750 ◦C, (**b**) initial powder sintered at 850 ◦C, (**c**) milled powder sintered at 750 ◦C and (**d**) milled powder sintered at 850 ◦C.

**Figure 5.** Examples of secondary electrons (SE) micrographs (**a**) initial powder sintered at 750 ◦C, (**b**) initial powder sintered at 850 ◦C, (**c**) milled powder sintered at 750 ◦C and (**d**) milled powder sintered at 850 ◦C.

*Metals* **2019**, *9*, 1280

The fraction of α-phase was determined by image analysis from SEM observations. Ten micrographs were analyzed for each sample. The results of this analysis are summarized in Figure 6. The conclusions from SEM observations were confirmed by image analysis. All sintered conditions contained a significant fraction of the α-phase, except for the initial powder sintered at 850 ◦C. As both sintering temperatures of 800 ◦C and 850 ◦C were above the reported temperature of β-transus (774 ± 14 ◦C) [1], α-phase should not precipitate during sintering. However, especially the milled powder was contaminated by oxygen, which is an α-stabilizing element and as such it increases the temperature of β-transus. This effect was studied in [31,33] both in Ti-Mo-Cr alloys and in pure Ti. It was found, that 0.4 wt. % of oxygen may increase the temperature of β-transus by 50 ◦C. The temperature of β-transus of milled powder can be therefore well above 800 ◦C. Therefore, in the initial powder sintered at 750 ◦C and in milled powder sintered at 750 ◦C and 800 ◦C, α-phase could precipitate during sintering. In initial powder sintered at 800 ◦C and 850 ◦C and in milled powder sintered at 850 ◦C, α-phase precipitated during cooling. High imposed strain in the milled powder provides preferential nucleation sites (dislocations and grain boundaries) for α-phase nucleation and enhances the diffusion promoting the growth of α-phase precipitates [34]. Both these effects—the higher contamination by oxygen and enhanced precipitation of α-phase due to milling—explain the higher amount of α-phase in the sintered milled powder compared to the initial one.

**Figure 6.** The evolution of α-phase content (vol. %) in sintered samples with sintering temperature.

The fraction of the α-phase decreased with increasing sintering temperatures for both sintered powders. The driving force for α-phase precipitation during sintering decreased as the temperature approached the temperature of β-transus. The amount of α-phase precipitated during cooling was clearly lower at higher temperatures.

The porosity of sintered both initial and milled powders depending on the sintering temperature is shown in Figure 7. Porosity of the initial powder decreased with increasing sintering temperature, while porosity of milled powder remained almost constant at all sintering temperature and was very low even at low sintering temperature. While high sintering temperature of 850 ◦C was necessary for sintering of the initial powder, the milled powder was well-sintered already at the lowest sintering temperature of 750 ◦C. It was caused mainly by different shapes of powder particles—the disk-shaped particles after milling are well stacked on each other and have also the higher surface area enhancing the efficiency of sintering as compared to ball particles in the initial powder. As a consequence, lower sintering temperatures were therefore sufficient to obtain fully compacted material. Moreover, the milled powder was better sintered due to enhanced diffusivity in the severely deformed and refined material [34].

**Figure 7.** Porosity of sintered initial and milled powder as a function of the sintering temperature.

The dependence of the microhardness on the sintering temperature for both initial and milled powder is shown in Figure 8. Microhardness was affected by residual porosity, by the contamination by oxygen and also by the content of both α-phase and ω-phase. While a decrease of porosity causes an increase in microhardness [19], a decrease in the oxygen content and in the amount of α-phase generally causes a decrease in microhardness [28,35]. The microhardness of sintered initial powder significantly increased with increasing sintering temperature, which was consistent with the decrease of the residual porosity. On the other hand, the amount of α-phase decreased with increasing sintering temperature. However, one might argue that the decrease in the amount of α-phase was compensated by an increase in amount of ω-phase. During cooling after sintering (cf. Figure 1), the Ti-15Mo alloy passes firstly through a region of α-phase precipitation in the temperature region of 750–500 ◦C [8,36]; and then through a region of ω-phase formation in the range of 450–250 ◦C [37]. One may therefore assume that ω-phase forms during cooling in the material sintered from the initial powder. On the other hand, the milled powder contained a considerable volume fraction of α-phase. In this case, the remaining β-matrix contained a higher amount of molybdenum due to element partitioning and was therefore comparatively more β-stabilized. Hence, the driving force for ω-formation was lower and the material with the higher amount of α-phase contained the lower amount of the ω-phase and vice versa.

This hypothesis was confirmed by XRD measurement of two conditions: samples sintered at 850 ◦C from the initial powder and the milled powder. Measured patterns are shown in Figure 9 along with LeBail fits, which allow only qualitative comparison of phase composition. Sintered initial powder contained considerable amount of ω-phase (highlighted by white arrows), while no ω phase was observed in the sintered milled powder. Milled condition contained substantial amount of α phase, which is consistent with SEM observations.

Sintered milled powder had a finer microstructure, higher contamination by oxygen, lower residual porosity and contained higher fraction of α-phase, in comparison with the sintered initial powder. All these effects contributed to the increased microhardness. However, as confirmed by XRD measurements, sintered milled powder did not contain the ω-phase. The microhardness of the sintered milled powder was therefore not enhanced in comparison with the initial powder (sintered at 850 ◦C) despite its refined microstructure. Coarsening of the microstructure (cf. Figure 5) and the decrease of the amount of α-phase with increasing sintering temperature (cf. Figure 6) did not affect microhardness of milled conditions. The dominant factor affecting microhardness might be increased oxygen content [38].

**Figure 8.** The temperature dependence of microhardness of sintered samples from initial and milled powder.

**Figure 9.** The x-ray diffraction (XRD) patterns of initial and milled powder sintered at 850 ◦C. Measured data are shown by black lines. LeBail fits are shown in color. White arrows point to the most pronounced peaks of ω-phase, black arrows highlight the most intensive α-phase peaks.

The maximum microhardness of both sintered powders was about 350 HV. It corresponds to the microhardness of α + β-structured Ti-15Mo prepared by conventional processing routes [35]. The microhardness of refined Ti-15Mo prepared by bulk SPD methods can be as high as 500 HV [28].

The results of tensile tests for selected samples are shown in Figure 10. For each sample, two specimens were deformed (noted as spec. 1 and spec. 2 in the Figure 10) and all successfully measured curves are shown. The flow curve for the milled powder sintered at 750 ◦C could not be shown, because the sample failed already in the elastic region. Milled powder sintered at 850 ◦C had comparable yield stress to the initial powder sintered at the same temperature. However, the initial powder sintered at 750 ◦C had lower yield strength. This was consistent with the microhardness measurement and could be explained the same way, namely by the concurrent effect of porosity, oxygen contamination, microstructure and phase composition. The yield strength of both powders sintered at 850 ◦C was around 1200 MPa. In comparison, coarse grained Ti-12Mo after two step ageing containing both β and α-phase possess yield strength of 1000 MPa [39]. Moreover, severely deformed metastable β-alloy TNTZO containing no α-phase had comparable yield strength over 1100 MPa [40].

**Figure 10.** Tensile tests: true stress–true strain curves.

Figure 11 shows a photo of tensile samples after testing. Sintered initial samples broke properly in the active part as marked by red (a) and green (b) arrows in Figure 11. On the other hand, none of the sintered milled powder samples broke in the middle of the active part, as marked by blue (c), yellow (d) and violet (e) arrows in Figure 11. This fact can be explained using the sketch in Figure 2. The cross section of the gauge length was only slightly smaller than the distance between the side of the sample and the holding pin hole as noted by the green line in Figure 2. This would not cause any problem in a standard isotropic material sintered from the initial powder. However, due to stacking of milled powder particles in the sintering die, the material was clearly stronger in the direction parallel to the tensile sample (i.e., perpendicular to the load during sintering). As the result, these samples failed near the pin hole because the flat powder particles were delaminated.

**Figure 11.** A photo of deformed tensile samples. Letters (**a**–**e**) correspond to the SEM images in Figure 12.

**Figure 12.** Fracture surfaces of selected specimens after tensile tests: (**a**) initial powder sintered at 750 ◦C; (**b**) initial powder sintered at 850 ◦C; (**c**) milled powder sintered at 750 ◦C; (**d**) milled powder sintered at 850 ◦C and (**e**) milled powder sintered at 850 ◦C-parallel to the powder particles.

This hypothesis was confirmed by SEM micrographs of fracture surfaces shown in Figure 12. The corresponding specimen and the position of SEM micrograph on the specimen is shown in Figure 11 by colored arrows. The fracture surface of initial powder sintered at 750 ◦C is shown in Figure 12a and corresponds to the region marked by the red arrow (a) in Figure 11. The initial round powder particles are clearly visible in Figure 12a. The sample was not fully sintered and the powder particles were joined only at the point of their contact.

A fracture surface of initial powder sintered at 850 ◦C corresponding to the zone marked by the green arrow (b) in Figure 11 is shown in Figure 12b. It shows a well sintered sample. No pores or initial powder particles were visible and the fracture was typical ductile. Figure 12c,d shows fracture surfaces of milled powder sintered at 750 ◦C and at 850 ◦C, respectively. They correspond to zones marked by blue and yellow arrows (c) and (d) in Figure 11, respectively. The fracture surfaces were completely different from those of sintered initial powder. Observed features were significantly finer in the Figure 12c (sintered at 750 ◦C) when compared to the Figure 12d (sintered at 850 ◦C), which corresponds well to the observed microstructures in Figure 5c,d. Highlighted regions by ellipses in Figure 12c,d corresponded to zones where a disk-shape powder particle was pulled out. Finally, Figure 12e shows a fracture surface of milled powder sintered at 850 ◦C oriented parallel to the powder particles as marked by a violet arrow (e) in Figure 11. In this SEM micrograph, planes corresponding to the surfaces of disk-shaped powder particles were visible. Milled powder particles were detached by the flat side of milled disk-shaped powder particles. These observations support our previous assumption of the possible delamination of powder particles.

Although well-sintered powders were ductile, the elongation of all sintered material was very poor, reaching maximum plastic strain of 3%. Commercial bulk Ti-15Mo alloy has a minimum elongation of 10% in aged α + β condition [1]. In similar Ti-12Mo alloy, the elongation of the α + β condition and the β condition was 9% and 50%, respectively [39]. Low ductility of sintered initial powder could be caused by the presence of pores and ω-phase, which results in the embrittlement of the material. Moreover, oxygen causes the embrittlement of α-phase. This phenomenon can affect the ductility especially in sintered milled powder, which is severely contaminated by oxygen and simultaneously contains high amount of α-phase.
