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For the SLM process, HIP and heat treatment decreased the yield strength and ultimate stress and improved the ductility. This can be ascertained from the test results and literature data [27,28,29,31,32]. From the publications reported, the machined and nonmachined specimens showed no obvious difference in terms of the ultimate tensile property; however, to improve the fatigue properties, the surface roughness needed to be machined to avoid surface defects such as open porosities easily becoming cracks when loaded. Therefore, all specimens in this work were surface-machined using the same standard. In addition, to verify the effect of HIP treatment on closing the internal defects, the density of the samples was measured, and the results are listed in Table 3.

LCF fractography of HT- and HIP-treated samples at different strain amplitudes. The dashed circles and dashed lines separate the crack growth regions of the fracture surface at different stages under corresponding strain. The strain amplitudes were 0.8%, 1.0%, 1.2%, and 2.0% (from left to right).

A fractography investigation of both types of SLM Ti-6Al-4V at four strain amplitudes was conducted to determine the effects of the heat treatment methods on crack initiation and propagation. It can be observed that the fracture features were similar at the specific strain amplitudes for both types of samples (Figure 14). For instance, fatigue striation can be found in the crack propagation area at lower strain amplitude (0.8%, 1.0%, and 1.2%), whereas it was scarce at the higher load level (2.0%). This suggests that the low-cycle fatigue of the samples exhibited different crack propagation performance at different strain amplitudes.

Low-cycle fatigue fracture surface and crack propagation features of HT- and HIP-treated samples: (a) HT at 0.8% strain amplitude; (b) HIP at 0.8% strain amplitude; (c) HT at 1.0% strain amplitude; (d) HIP at 1.0% strain amplitude; (e) HT at 1.2% strain amplitude; (f) HIP at 1.2% strain amplitude; (g) HT at 2.0% strain amplitude; (h) HIP at 2.0% strain amplitude.

In addition, it was found that there was more than one crack initiation location for both types of samples, and no defect was found close to the crack initiation. A typical fatigue fracture surface of HIP samples at 1.0% strain amplitude was investigated (Figure 15). Similar to typical fatigue fractography, the crack initiated from the surface, and striations could be found in the crack growth area (Figure 15a,b); the propagation steps were also found (Figure 15c,d).

Typical low-cycle fatigue fracture surface features of HIP-treated sample at 1.0% strain amplitude: (a) morphology of a low-cycle fatigue specimen (FCG denotes crack propagation); (b) crack initiation region; (c) striations in the crack growth region; (d) transition region with striations and steps.

Schematic of internal defects along the fatigue crack propagation of SLM HT and HIP samples: (a) HT samples with pore defects; (b) internal defects on the fracture surface; (c) crack propagation of HIP sample; (d) typical region of fracture surface.

In general, the fatigue surface includes the fatigue crack initiation and propagation regions, as well as the final shear fracture region. In this research, rather than relying on initiation from pores, it was found that cracks initiated at alpha colony boundaries in the region close to the outer surface of the samples. At the initiation stage, microcracks prefer propagating along α-laths within the mixed (αs+β) microstructure [41,47]. HT or HIP α+β titanium alloys with lamellar microstructures have similar crack initiation mechanisms. In addition, as reported in [48], cracks in AM parts are normally generated from α colonies with larger grain sizes. Compared to HT samples, microcracks were easily originated on the HIP sample surface. The crack initiation sites and crack propagation areas are shown in Figure 20. Figure 20a (HT) and Figure 20d (HIP) show the three-dimensional scanned fracture surfaces, which could be divided into three areas. The first is the crack initiation area, while the second is the propagation region. Some ridges due to the effect of high stress can be observed on the fracture plane (shown in Figure 20a,d). Compared to the HIP sample, numerous paralleled branch secondary cracks could be observed on the HT sample fracture plane. This phenomenon can be regarded as one reason why HT samples had a superior LCF life to HIP samples, because the energy was dispersed to multibranch cracks, and the crack propagation trend could be retardant [29].

Fracture surfaces and crack propagation path: (a) fracture surface scanning model of HT sample at 0.8% strain amplitude; (b) crack initiation region of HT sample; (c) crack propagation path of HT sample; (d) fracture surface scanning model of HIP sample at 0.8% strain amplitude; (e) crack initiation region of HIP sample; (f) crack propagation path of HIP sample.

With the increase in crack propagation, the fatigue cracks can penetrate α-laths; the intragranular crack propagation can be observed in Figure 20c (HT) and Figure 20f (HIP). Cracks that propagated along the α + β grain lath boundaries can be also observed in the HT crack path (Figure 20c). It is shown that crack tip turned and was deflected more frequently within the finer microstructure of the HT sample, leading to a zig-zag crack path and a rougher fracture surface on the micro and mesoscale.

where NInc is the number of cycles to incubate a crack, NMSC is the number of cycles for the growth of an MSC, NPSC represents the life cycles for the growth of a PSC, and NLC represents the life cycles for long crack growth to final failure. For SLM Ti-6Al-4V, NLC is not observed and can be neglected [22,23].

As shown in Figure 21, the fatigue life prediction results of the two heat-treatment processes obtained using the MSF model were plotted. Figure 21a,c exhibit the total fatigue life and the fractions, incubation life, and MSC/PSC life of HT and HIP samples, respectively. Figure 21b,d present the upper and lower bounds of the fatigue life according to experimental data; here, we assumed that the cracks formed and initiated at casting pores of diameter 5 and 500 µm, respectively. Taking the dispersity and the estimation of parameters in the MSF model into consideration, the prediction fatigue life curves showed agreement with the experiment results. However, at a high strain amplitude level (Δεa/2 = 2%), there was a noticeable error between the experimental data and prediction. This is because only one sample was evaluated at high strain amplitude, leading to more unpredictable low-cycle fatigue results. Therefore, the MSF prediction model was deemed acceptable in this research. The upper bound and prediction life adequately fit the experiment data, as we assumed the smallest and largest incubation pore diameters as 5 and 500 µm; thus, actual internal defects or voids of size smaller than 80 µm could have similar effects on the prediction results (Table 7).

Fatigue life prediction using MSF model for HT and HIP samples: (a) crack incubation and small crack life for HT specimens; (b) lower bound and upper bound for HT specimens in MSF model; (c) crack incubation and small crack life for HIP-treated specimens; (d) lower bound and upper bound for HIP-treated specimens in MSF model.

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