*3.2. Mechanical Properties*

Figure 6 shows the hardness (normal direction) comparison of AS, 400-AC, 600-AC, and 800-AC. The hardness of AS, 400-AC, and 600-AC is similar, but when the heat treatment temperature increases up to 800 ◦C, the hardness decreases because α' phases are completely transformed into α + β phases.

**Figure 6.** Hardness of AS, 400-AC, 600-AC, and 800-AC (normal direction).

The room temperature tensile properties of AS, 400-AC, 600-AC, and 800-AC are shown in Figure 7, with the mean value of the mechanical properties list in Table 4. It can be seen that the strength of 400-AC is significantly higher than that of AS. According to previous reports [23,24], the tensile strength can be improved by the precipitation of α phases in the grains or on the boundaries, and the residual stress is also reduced by a 400 ◦C heat treatment, so the strength of 400-AC is higher than that of AS.

**Figure 7.** Tensile properties of AS, 400-AC, 600-AC, and 800-AC at room temperature for (**a**) strength and (**b**) ductility.

**Table 4.** Room temperature tensile properties of AS, 400-AC, 600-AC, and 800-AC. (YS: yield strength; UTS: ultimate tensile strength; UE: uniform elongation; TE: total elongation).


The β phases begin to precipitate, and the strength gradually decreases with increases in the heat treatment temperature, and the strength is similar to AS at 600 ◦C, but the strength at 800 ◦C is lower than that of AS because the α' phases are completely transformed into the α + β phases. On the other hand, the ductility increases as the heat treatment temperature increases, but is still less than 10% and less than that of AS. The precipitation of the α phases increases the strength but significantly decreases

the ductility. Increases in the ductility are attributed to the formation of β phases, and α' phases are completely transformed into α + β phases upon increases in the heat treatment temperature. In other words, the precipitation of the α phases contributes to the improvement of tensile strength, where the increase in ductility can be attributed to the formation of the β phases.

Figure 8 shows the morphology of the tensile fracture surfaces of AS, 400-AC, 600-AC, and 800-AC at room temperature. Dimpled ductility structures can be observed in all specimens. In addition, some cleavage facets can be observed in 400-AC, which is consistent with its low ductility.

**Figure 8.** Morphologies of tensile fracture surfaces at room temperature for (**a**) AS, (**b**) 400-AC, (**c**) 600-AC, and (**d**) 800-AC.

Figure 9 shows the high temperature tensile properties of AS, and the mean value of the mechanical properties is listed in Table 5. The strength decreases slowly as the temperature increases. The UTS is close to 1000 MPa and the YS is about 800 MPa at 400 ◦C, indicating that the phase transformation α'→ α + β is exhibiting decreased strength. Compared to previous reports [25,26], the strength of commercial Ti-6Al-4V has a dramatic decrease whilst the temperature increases. Apparently the as-SLM Ti-6Al-4V could maintain a certain strength under 400 ◦C. In terms of ductility, there is first an increasing and then a decreasing trend, with the total elongation (TE) close to 10% at 250 ◦C–300 ◦C, which means the 3D-printer titanium alloy can be used in medium temperature applications. It is worth noting that the ductility at 400 ◦C is similar to that of AS, which implies that it is affected by a small α phase precipitation effect. In summary, the tensile properties of 350 ◦C are most applicable with the highest values in uniform elongation (UE). Figure 10 shows macroscopic morphology photographs of high temperature tensile fracture specimens, though there is not an obvious shrinkage phenomenon in any of the specimens, and there is no hole expansion at the location of the inserted pin. The macroscopic fracture morphology becomes gradually flatter from a zigzag pattern as the tensile temperature increases.

**Figure 9.** High temperature tensile properties of AS for (**a**) strength and (**b**) ductility.


**Table 5.** High temperature tensile properties of AS.

**Figure 10.** Macroscopic morphology photographs of room temperature and high temperature tensile fracture specimens.

#### *3.3. Particle Erosion Characteristics and Mechanisms*

In order to clarify the difference in erosion characteristics of different matrices, this study chose AS and 800-AC specimens exhibiting erosion for comparison. Figure 11 shows the erosion data of AS and 800-AC eroded by Al2O3 particles; there are at least three specimens for each test and the mean value of the test specimens was taken as the erosion results of the corresponding condition. The reason for using AS for comparison with 800-AC was due to the fact that the martensitic α' phases of AS were completely transformed into the lamellar α + β phases, where the ratio α:β was about 68:32. Because AS and 800-AC exhibited completely different microstructures, the effects of the phase difference on the erosion wear properties could be fully investigated. Notably, both AS and 800-AC had maximum erosion rates at 30◦ impact, where the erosion rates decreased with increases in the impact angle, and the minimum erosion rate at a 90◦ impact resulted in ductile-cutting dominating the erosion behavior. At all impact angles, the erosion rate of AS was greater than that of 800-AC, indicating the erosion resistance of the continuous lamellar α + β phases of 800-AC were better than that of AS. The erosion resistance was positively correlated with the hardness of conventional Ti-6Al-4V alloy [27], but the result was different in this research. The difference in microstructure caused this result, therefore we investigated the surface and subsurface morphologies of AS and 800-AC at different impact angles.

**Figure 11.** The erosion rate as a function of the impact angles of AS and 800-AC.

Figure 12a shows the surface morphology of AS at 30◦ impact, where lips and grooving can be seen on the surface. Figure 12b shows the surface morphology of AS at 90◦ impact, where pits formed by erosion wear can be seen. Figure 12c shows the surface morphology of 800-AC at 30◦ impact, where compared with the lips on the AS surface, the surface morphology of 800-AC was completely cut by erosion particles and the lips of the erosion subsurface were smaller. Figure 12d shows the surface morphology of 800-AC at 90◦ impact, where pits formed by erosion wear can also be seen, similar to AS. However, some scratch marks could also be observed on the surface of 800-AC at 90◦ impact. This phenomenon indicated that the specimen was deformed by squeezing first and then was stripped by the erosion particles.

**Figure 12.** Surface morphologies of AS at (**a**) 30◦ and (**b**) 90◦. Surface morphologies of 800-AC at (**c**) 30◦ and (**d**) 90◦.

Figure 13a shows subsurface morphologies of AS at 30◦ impact, where lips caused by erosion particles can be found on the subsurface in the circled area. Figure 13b shows subsurface morphologies of AS at 90◦ impact, where both narrow and deep pits can be observed, and subsurface morphologies of 800-AC at 30◦ impact, where almost no lips can be observed (Figure 13c). They are replaced by smoother features resulting from erosion particles. Figure 13d shows the subsurface morphologies of 800-AC at 90◦ impact, where the pits that can be observed in 800-AC are wider and shallower than those in AS. Notably, the surface roughness of erosion subsurface is lower than that of AS-90◦ impact.

**Figure 13.** Subsurface morphologies of AS at (**a**) 30◦ and (**b**) 90◦. Surface morphologies of 800-AC at (**c**) 30◦ and (**d**) 90◦.

Figure 14 provides the erosion mechanism schematic diagrams of AS and 800-AC at impacts of 30◦ and 90◦, respectively. Figure 14a is the erosion mechanism of AS at 30◦ impact, where it can be seen that due to the inability of the needle-like martensitic α' phases to prevent erosion wear due to erosion particles, the lips are formed by the cutting mechanism. When AS is eroded by erosion particles at 90◦ impact, the surface of the specimen is subjected to the positive impact of erosion particles, which form squeeze features on the erosion wear surface, but the positive impact resistance of AS is weaker than that of 800-AC, thus resulting in the formation of deeper pits, as shown in Figure 14b. When the 800-AC is inclined and eroded by erosion particles at 30◦ impact, the lamellar α + β phases process smaller and fewer lips, as shown in Figure 14c. Figure 14d shows the erosion wear behavior of 800-AC at 90◦ impact, where some of the continuous lamellar α + β phases are scraped off when the erosion particles have a positive impact. However, the damage characteristics are not significant.

**Figure 14.** Particle erosion mechanism of AS at (**a**) 30◦ and (**b**) 90◦. Particle erosion mechanism of 800-AC at (**c**) 30◦ and (**d**) 90◦.

According to the discussion above regarding the eroded surface morphology, including the subsurface morphology of the erosion and the erosion mechanism, although the ductility of 800-AC was not as good as that of AS, its continuous lamellar α + β phases were tough enough to limit the ability of the erosion particles to scrape off the material at all impact angles. Thus, the erosion resistance of 800-AC was determined to be better than that of AS. The SLM Ti-6Al-4V specimen must be heat treated to have wear resistance.
