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Article

Effects of HIP Process Parameters on Microstructure and Mechanical Properties of Ti-6Al-4V Fabricated by SLM

by
Zhoujin Lv
1,
Haofeng Li
1,
Lida Che
1,
Shuo Chen
1,*,
Pengjie Zhang
1,
Jing He
1,
Zhanfang Wu
1,
Shanting Niu
1 and
Xiangyang Li
2,*
1
CISRI HIPEX Technology Co., Ltd., Beijing 100081, China
2
Beijing Advanced Innovation Center for Materials Genome Engineering, Central Iron & Steel Research Institute, Beijing 100081, China
*
Authors to whom correspondence should be addressed.
Metals 2023, 13(5), 991; https://doi.org/10.3390/met13050991
Submission received: 26 April 2023 / Revised: 12 May 2023 / Accepted: 15 May 2023 / Published: 20 May 2023
(This article belongs to the Special Issue Titanium Alloys: Microstructure, Deformation and Mechanical Properties)
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Versions Notes

Abstract

:
Ti-6Al-4V titanium alloy products formed by selective laser melting (SLM) are characterized by high strength and low plasticity. In addition, there may be pores inside the material, which may become a fracture sprouting point and accelerate the failure of the parts. Using an optical microscope (OM), scanning electron microscope (SEM), and electronic universal testing machine, the effects of hot isostatic pressing (HIP) parameters on the microstructure and tensile property of SLM-formed Ti-6Al-4V titanium alloy were investigated. The results show that HIP performed below the β-phase transition temperature, and the structure of the Ti-6Al-4V titanium alloy is composed of an α phase and β phase. With the increase in the HIP temperature, the α lath coarsens into a short rod, the content of the β phase increases and coarsens, and the tensile strength and yield strength of Ti-6Al-4V show a decreasing trend. With an HIP process performed at a temperature of 910 °C and pressure of 130 MPa for 2 h, the Ti-6Al-4V titanium alloy obtains the best matching of strength and plasticity.

1. Introduction

Selective laser melting (SLM) is one of the most important branches of additive manufacturing (AM) techniques. During the last few years, SLM has attracted much attention due to its inherent advantages, such as high material utilization, short cycle times, and insensitivity to geometric shapes [1,2,3]. Ti-6Al-4V, an α + β-type titanium alloy, is the most widely used titanium alloy, which makes up roughly half of all titanium alloys used commercially. This material possesses excellent corrosion resistance, high-temperature mechanical properties, and biocompatibility; it is widely used in aerospace, petrochemical, and biomedical engineering [4,5,6]. Nevertheless, the low thermal conductivity, propensity for strain hardening, and chemical reactivity to the oxygen of Ti-6Al-4V make manufacture extremely challenging. Therefore, SLM is recognized as a significantly more efficient approach to fabricating complex-shaped Ti-6Al-4V parts [1,2]. Due to the higher temperature gradient involved in SLM, the microstructure of Ti-6Al-4V alloy processed by SLM is α′ acicular martensitic, and due to epitaxial growth, elongated grains emerge [7,8]. Thus, SLM components made of Ti-6Al-4V generally possess high tensile strength but poor ductility. Due to its inherent enormous temperature gradient, SLM workpieces are prone to significant residual stresses, which lead to a tendency for the workpieces to crack and deform [1]. In addition, there may be pores within the material, which often become the starting point for cracking when the workpiece is subjected to load, accelerating the fatigue failure of the products [3].
Given the above problem, experts and scholars have made a lot of exploration of Ti-6Al-4V powder raw materials [9,10], the SLM forming process [11,12,13], and the subsequent heat treatment process [14,15]. Ordinary heat treatment cannot eliminate internal defects such as pores [16]. Hot isostatic pressing (HIP) technology originated in the 1950s, first for diffusion joining of fuel elements in the atomic energy reaction process [17]. Due to its own process superiority, HIP was gradually expanded to various production areas such as castings densification treatment, diffusion connection of heterogeneous materials, and powder metallurgy near-net forming [17,18,19]. As an attractive post-treatment technology, the high temperature and pressure of the HIP process can significantly improve the structure of the workpiece, eliminate internal porosity, and improve the overall mechanical properties of the material [20,21,22]. Currently, important castings applications in aerospace and the military are required in HIP densification treatment [17]. HIP technology is also used in the post-processing of additively manufactured parts, and the combination of 3D printing technology and isostatic pressing technology has become a hot topic of research, providing a convenient implementation solution for the isostatic pressing of structurally complex products [19,23,24,25,26,27]. The effects of HIP on SLM alloys such as aluminum alloy [23], magnesium alloy [19], titanium alloy [24,25], nickel alloys [26], and steels [27] have received significant attention. Hot isostatic pressing is a combination of high temperature and pressure. The introduction of pressure affects the material’s phase changes and interfacial reactions, and the parameters of the hot isostatic pressing process play a crucial role in the material properties. However, most of the experts and scholars investigating the influence of the HIP treatment process on Ti-6Al-4V titanium alloy after SLM forming have only studied one HIP treatment process, or the parameters such as temperature and pressure vary drastically between different HIP processes [2,3,4,6,10,12,16]. Within a narrower process range, the effect of systematic HIP treatment processes on the properties of SLM-formed Ti-6Al-4V titanium alloys is still less studied. Based on this understanding, this paper investigates the effects of different hot isostatic pressing temperatures and pressure combinations on the microstructure and tensile properties of Ti-6Al-4V titanium alloy formed by SLM. The aim is to provide a foundation for the control of microstructure and optimization of mechanical properties during the HIP treatment of SLM-formed Ti-6Al-4V titanium alloy parts.

2. Materials and Methods

Ti-6Al-4V titanium alloy powder prepared by electrode induction melting gas atomization technology (HEIGA-S, BLT, Xi An, China) was used as SLM forming material, the chemical composition of the powder is shown in Table 1, and the powder’s physical properties’ results are shown in Table 2. BLT-S310 was used as the forming machine (BLT, Xi An, China) max. forming specimen size: 250 mm × 250 mm × 400 mm, laser power 500 W). Argon gas with 99.99% purity was used as the protective gas. Specimens were built along the Z-axis, and the forming specimens’ sizes were ϕ12 mm × 65 mm. The detailed printing parameters are shown in Table 3.
ASTM F2924-14 specifies that HIP is required for Class C components and optional for all other classifications produced by powder bed fusion [28]. Components are processed under an inert atmosphere at no less than 100 MPa within the range 895 °C to 955 °C, held at the selected temperature within ±15 °C for 180 ± 60 min, and then cooled under an inert atmosphere to below 425 °C [28]. In order to study the effect of the holding temperature and pressure of the HIP process on the microstructure and mechanical properties of Ti-6Al-4V titanium alloy formed by SLM, 9 different HIP schedules were designed for this experiment. For each HIP process, the applied temperature and pressure are shown in Table 4. All the HIP processes maintain a holding time of 2 h, a heating rate of about 5 °C/min, and a cooling rate of about 3 °C/min.
Using Archimedes’ method, the specimens were immersed in deionized water at 25 °C to determine their density. Tensile specimens were obtained by cutting from SLM-formed and HIP-treated cylindrical samples using a computer numerical control lathe (TK36S, Baoji Machine Tool Group, Baoji, China). The geometric parameters of the tensile specimens are listed in Figure 1. The tensile properties of the Ti-6Al-4V titanium alloy at room temperature were performed on a microcomputer-controlled electronic universal testing machine. According to ISO 6892-1, before the specimens yielded, the stretching speed was 0.5 mm/min, and after the specimens yielded, the stretching speed was 5 mm/min. Each process measured 5 tensile specimens. By tungsten scanning electron microscopy (SEM, EVO25, ZEISS, Oberkochen, Germany), the fracture morphology of the tensile specimens was observed. Metallographic specimens were prepared by a wire cut electric discharge digtal control machine tool (DK7732, Taizhou, China) cutting cross-sections and longitudinal sections of the cylindrical samples in different processing conditions. Before microstructure observation, all specimens were ground, polished, and etched with Keller reagent (HF:HNO3:H2O = 1:2:17). Metallographic microscopy (OM, DMI8A, Leica, Weztlar, Germany) and scanning electron microscopy (SEM, Quanta 650 FEG, FEI, Hillsboro, America) were used to observe the microstructure morphology of the specimen before and after hot isostatic pressing. ImageJ software was used to analyze the α-lath thickness or interlamellar spacing λα of HIP-treated specimens. Drawing on Vander Voort and Roosz’s [29] method for measuring the interlayer spacing of pearlite in AISI 1040 carbon steel samples. Haize Galarraga [30,31] also applied this method to measure the α-lath thickness of Ti-6Al-4V ELI in different heat treatment states after EBM forming. The method consists of calculating the mean random spacing, σr, using the intersection method and then estimating the mean true spacing, σt. Draw a circle with a radius of R on the high magnification metallographic pictures in an unbiased manner (Figure 2). The number n of intersections of the α lamellae with the test line is counted. The number of intersections per unit length of the test line, NL, is calculated by dividing n by the true total test line length per measurement field. The reciprocal of NL is the mean random spacing, σr. Finally, the mean true spacing, σt, is calculated by dividing the mean random spacing by two. This methodology has been applied five times in one image. Three images were selected for each HIP schedule. These values were averaged to determine the interlamellar spacing at each schedule. ImageJ software (ImageJ 1.53c) is also used for measuring β-phase volume fraction, Vβ. The image to be analyzed was imported into the software. An image with a clear contrast between α phase and β phase was obtained by utilizing Image-Adjust-Threshold adjustment. The percentage of area occupied by the α phase was calculated by the software. The percentage content of β phase can be obtained by subtracting the percentage area of α phase from 100%.

3. Results

3.1. Microstructure

The microstructure of the Ti-6Al-4V titanium alloy specimen after SLM forming is shown in Figure 3. Figure 3 exhibits a typical microstructure of SLM-formed Ti-6Al-4V. Clear prior β-grain boundaries can be observed in both cross-section and longitudinal sections. The cross-sectional microstructure consists of equiaxed prior β grains with a grain size of approximately 110 μm. The microstructure of the longitudinal section consists of columnar primary β grains with epitaxial growth. The width of the columnar crystals is comparable to the grain size of the equiaxed crystals in the cross-section. Kumar et al. [15] suggested that this size is determined by the hatch spacing. Lore Thijs et al. [32] have analyzed the causes of this microscopic morphology. In the SLM forming process, the temperature gradient of the molten pool is perpendicular to the laser scanning surface. The bottom of the molten pool is where the molten pool begins to melt. As the laser beam scans the powder layer, it remelts the tip of the previous layer of columnar crystals. Thus, the top of the unfused columnar crystal becomes the directionally solidified nucleus of the layer. Therefore, the original columnar grains will continue to grow epitaxial along the deposition direction [32].
The basket weave microstructure of fine acicular α′ martensite phases staggered distribution with various orientations inside the grains within the primary β grains. During the SLM process, the cooling rate of the high-temperature β phase is exceptionally rapid, about 104~106 K/s, exceeding the critical martensite cooling transition temperature [33]. In Ti-6Al-4V titanium alloy, the Burgers orientation relationship (BOR) exists between the low-temperature hexagonal close-packed (bcc) α phase and the high-temperature body-centered cubic (bcc) β phase [15,34]. It is based on two parallel planes (0001)α//{110}β and directions <1120>α//<111>β in the respective phases. Due to the symmetry of the parent and product crystals, there are 12 different variants of the α/α′ nucleation of the β matrix [15,35]. The extreme cooling rate during the SLM forming process provides a sufficient driving force for all the variants that can be formed, ultimately leading to the formation of multiple α′ martensitic phases with various orientations [15,35]. According to the SEM images in Figure 3c,d, it can be observed that the SLM-formed Ti-6Al-4V is partially distributed with white nano dot clusters inside the coarser acicular α′ martensite phase. According to Ail et al. [36], such nano dot clusters are nanoscale β phases. The nano-β phase is nucleated between the incipient martensite in a pure Ar gas environment during the SLM forming process.
Figure 4 shows the cross-section microstructure of Ti-6Al-4V specimens treated by different HIP schedules. The microstructure shows a net basket structure consisting of an α phase (light color) and β phase (dark color). The effect of HIP temperature on the microstructure is more obvious compared to the HIP pressure. At the same pressure, the HIP temperature was changed by only 20 °C each time; however, delicate microstructural differences can still be observed. In Figure 4a–c, after hot isostatic pressing treatment at 910 °C, a majority of the α phase is a lathy structure, and a few of the α phases show a rod-like structure. In Figure 4d–f, after HIP treatment at 930 °C, the α phase coarsens, and the rod-like α phase increases. In Figure 4g–i, the spheroidized and coarsened α phase is further increased after HIP treatment at 950 °C. In Figure 4i, showing HIP treatment with T3P3, the width of the α phase in the Ti-6Al-4V titanium alloy specimen has exceeded 5 μm.
To more accurately characterize the effect of the HIP schedules on the α-lath dimensions, Figure 5 shows the statistics of the α-lath widths of the specimens’ cross-sections after different HIP processes were calculated by ImageJ software. It can be seen through Figure 5, at the same HIP pressure, that the average width of α laths gradually increases with the increase in HIP temperature. At the same HIP temperature, the average width of α laths showed a trend of first decreasing and then increasing with the increase in HIP pressure. As the HIP schedule of T1P2, the average width of α laths was about 0.918 μm, which was the smallest among all HIP schedules. As the HIP schedule of T3P3, the average width of α laths was about 1.49 μm, which was the largest among all HIP treatment schedules.
In the cross-sectional SEM photographs of Ti-6Al-4V specimens after different HIP processes, as can be seen in Figure 6, the β phases exhibit two different morphologies. One is the irregular micron-scale β phase formed at the boundary of the α phase, and the other is the nano-dotted β phase diffusely distributed inside the α-lath structure. By HIP treatment at different temperatures, the variation of irregular β shows the same trend as the variation of the α phase. As the HIP treatment temperature increases, the α laths coarsen; similarly, the β phases grow up. It can be observed from Figure 6i that the equivalent diameter of the larger size β phase exceeded 3 μm following the HIP treatment at T3P3.
The β-phase volume fractions (Vβ) in the Ti-6Al-4V specimens with different HIP schedules are shown in Figure 7. The effects of HIP temperature and HIP pressure on Vβ follow the same trend. With the increase in HIP temperature, the content of the β phase gradually increases. Similarly, the content of the β phase gradually increases with the increase in HIP pressure. The Vβ of the HIP-treated specimens at 950 °C was significantly higher than that of the HIP-treated specimens at 910 °C and 930 °C. After the HIP treatment at T1P1, the Vβ was only 13.07%. Following the HIP treatment at T3P3, the Vβ increased to 27.85%.
X-ray energy spectroscopy (EDX) analysis was performed on the HIP-treated specimen at T2P3 to help with identifying the phase composition. The result of EDX analysis is shown in Figure 8. Al is an α-phase stable element with a high content in the α phase, and V is a β-phase stable element with a higher content in the β phase [37,38]. The EDX results also show that the content of the Ti element in the α phase is higher than that in the β phase.

3.2. Physical Properties

The density test results of specimens are shown in Figure 9. The theoretical density of Ti-6Al-4V titanium alloy is about 4.43 g/cm3 [2]. The density of SLM-formed Ti-6Al-4V reached 4.394 g/cm3, which is equivalent to 99.2% of the theoretical density. After HIP treatment, the density of all samples was above 4.415 g/cm3, which exceeded 99.7% of the theoretical density. HIP treatment can heal the internal microporosity defects and further improve the densities of SLM-formed Ti-6Al-4V.
Figure 10 shows the tensile properties of SLM-formed Ti-6Al-4V specimens following different HIP processes. The tensile properties of all specimens exceeded the requirements in the ASTM F2924-14 standard [28]. The following are the key observations made from the data given in Figure 9: (I) At the same HIP pressure, the tensile strength (UTS) and yield strength (YS) of specimens decrease with increasing HIP temperature. (II) At the same HIP temperature, the tensile strength and yield strength of SLM-formed Ti-6Al-4V first increase and then decrease with increasing HIP pressure. After the HIP treatment process of holding at 910 °C and 130 MPa for 2 h, the specimens obtained the optimal static tensile properties in this study. The average YS of the specimens was 935 MPa, the average UTS was 1005 MPa, and the average elongation (EL.) was 16.5%.
To compare the relationship between microstructure and mechanical properties, the information on tensile properties, λα, and Vβ treated with the different HIPs is summarized in Table 5. Beyond HIP treatment, both YS and UTS decrease in all cases. The overall range of YS reductions is between 15% and 20%, and the overall range of UTS reductions is between 18% and 23%. While YS ranges between 883 MPa and 935 MPa, UTS varies from 955 MPa to 1005 MPa. These strength values are still superior to those of the Ti-6Al-4V titanium alloy formed by conventional manufacturing processes [1,39,40]. The difference between UTS and YS shows that the amount of strain hardening is relatively minor for all samples. The SLM-formed specimens’ yield ratio reached 90%, and following HIP treatment, the yield ratio of the Ti-6Al-4V further increased to 92~99%. It should be noted that the elongation of SLM-formed specimens is reaching 10.3%, which meets the requirements of the ASTM F2924-14 standard for the elongation of Ti-6Al-4V titanium alloy, and this experiment is a further enhancement of 38.8~60.2% based on SLM-formed specimens.
The fracture morphology of tensile specimens is shown in Figure 11. The macroscopic fracture morphology of all the specimens showed a cup–cone fracture surface. The SLM-formed specimen fracture surface consisted of a fiber zone and a larger proportion of the peripheral shear lip zone, in which macroscopic fracture necking was not obvious. Fractures sprout in the center of the section and then expand into the surrounding area. The microscopic fracture morphology shows a mixed fracture mechanism of river-like crystal penetration decomposition and coexisting dimples. The dimples are shallow and sparse. After HIP treatment, the macroscopic fractures of Ti-6Al-4V specimens all exhibit a cup–cone fracture surface with significant striction and a smaller proportion of peripheral shear lip. The microscopic fracture morphology is characterized by deep and dense equiaxial dimples, which demonstrate typical ductile fracture characteristics and indicate that the materials have good ductility.

4. Discussion

The unique thermal cycling history of the SLM fabrication process causes four different sizes and morphologies of martensitic phases that can be formed in Ti-6Al-4V [41]. HIP treatment of SLM-formed Ti-6Al-4V in the two-phase region transformed the sub-stable α′ martensitic phase into a mixed α + β two-phase microstructure [16]. During the high-temperature treatment, the α′ martensite phase decomposes, the α phase nucleates along the α′ martensite phase boundary, and the V atoms are expelled, resulting in the inhomogeneous nucleation of β at the α-phase boundary [42]. Coarsening of the α laths and growth of the β phase also occur at high-temperature treatment below the β-phase transition temperature, and the net-basket-shaped α laths gradually transform into short rods [43]. As the HIP temperature approximately approaches the β-phase transition temperature, both α and β phases’ growth of grains become more pronounced [15]. Compared to the HIP temperature, the effect of HIP pressure on the grain size and phase transition of SLM-formed Ti-6Al-4V is more complex. For a relatively lower HIP temperature of 910 °C, the influence of pressure on grain size and the β-phase transition is minor. Under the higher HIP temperature of 950 °C, a more significant influence on grain size and the β-phase transition was observed. This study was conducted based on the ASTM F2924-14 standard, and the results showed that the lower temperature (910 °C) and appropriate pressure (130 MPa) helped to suppress the grain growth, and the effect of a more extensive temperature and pressure on the grain size and phase transition of Ti-6Al-4V remains to be investigated in depth.
The tensile properties of SLM-formed Ti-6Al-4V are mainly dependent on the internal defects of the material as well as the content and morphology of α′, α, and β phases [1,13,16]. The α′ martensite in SLM-formed Ti-6Al-4V provides superior strength and inferior plasticity [14,37]. In addition, a high density of dislocations within α′ martensite prevents dislocation migration. Both mentioned facts combined to enhance the strength of SLM-formed Ti-6Al-4V [15]. As expected, the tensile strength and yield strength of Ti-6Al-4V decreased following the HIP treatment, but the elongation increased significantly. With HIP treatment below the β-phase transition temperature, the tensile and yield strengths of SLM-formed Ti-6Al-4V decrease following the HIP temperature. α′ martensite decomposition into the α + β two-phase microstructure during HIP resulted in a decreased strength and increased plasticity of the material [16]. A transition from α to β phase and coarsening in both the α and β phases also occur during the HIP treatment [16]. HIP at 910 °C, fine and numerous α-phase grains are observed within the Ti-6Al-4V titanium alloy, and more slip systems could be driving in α phase. While the specimens undergo plastic deformation, the slips are first moved in individual α grains, and then the deformation spreads to other grains immediately, so it is difficult to attract stress concentration and fracture in individual grains. Therefore, the microstructure of the α phase with finer and numerous grains exhibits more deformation and better plasticity [15,16]. An hcp structure α phase has only three slip systems, while a bcc structure β phase has 12 slip systems. Along with the increase in the HIP temperature, on the one hand, the α-phase quantity decreasing follows the β-phase quantity increasing; thus, the material has improved plasticity. On the other hand, the α phase swallows up and expands among each other, which causes a decrease in both strength and plasticity of the titanium alloy [15,43].

5. Conclusions

This study investigated the effects of HIP treatment temperature and pressure within the standard requirements of ASTM F2924-14 on the microstructure and mechanical properties of SLM-formed Ti-6Al-4V titanium alloy. The content and morphology of the α phase and β phase in the Ti-6Al-4V using various HIP process treatments were characterized. The following conclusions can be drawn from this study:
  • At the same HIP pressure, by increasing the HIP temperature, the percentage content of the β phase gradually increases, and the α and β phases swallow each other and grow up, but the quantity decreases. At the same HIP pressure, by increasing the HIP temperature, the β-phase percentage increases gradually, and the α-lath width tends to decline first and then rise.
  • At the same HIP pressure, the tensile strength and yield strength of SLM-formed Ti-6Al-4V decline with elevated HIP temperature. At the same HIP temperature, the tensile strength and yield strength of SLM-formed Ti-6Al-4V rise first and then decline following HIP pressure.
  • With an HIP process temperature of 910 °C and pressure of 130 MPa for 2 h, the α-lath width in SLM-formed Ti-6Al-4V was about 0.956 μm, while the β-phase bulk percentage content was 15.07%. The material obtained an optimal strength and ductility match with a tensile strength of 1005 MPa, yield strength of 935 MPa, and elongation up to 16.5%.

Author Contributions

Conceptualization, Z.L., S.C., and X.L.; methodology, Z.L., H.L., and P.Z.; software, L.C.; validation, Z.L. and S.N.; formal analysis, Z.L. and Z.W.; investigation, Z.L. and L.C.; resources, Z.L.; data curation, Z.L., H.L., and J.H.; writing—original draft preparation, Z.L.; writing—review and editing, X.L., S.C., and L.C.; visualization, Z.L.; supervision, S.C.; project administration, Z.L.; funding acquisition, Z.L. All authors have read and agreed to the published version of the manuscript.

Funding

This research was funded by Beijing Nova Program (No. 2022139).

Institutional Review Board Statement

Not applicable.

Informed Consent Statement

Informed consent was obtained from all subjects involved in the study.

Data Availability Statement

The data presented in this study are available in the article.

Conflicts of Interest

The authors declare no conflict of interest.

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Figure 1. Tensile test specimen geometry.
Figure 1. Tensile test specimen geometry.
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Figure 2. The intersection of α laths on a 9 μm radius circle was measured with the ImageJ software.
Figure 2. The intersection of α laths on a 9 μm radius circle was measured with the ImageJ software.
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Figure 3. Microstructure of SLM specimen: (a,c) cross-section; (b,d) longitudinal section.
Figure 3. Microstructure of SLM specimen: (a,c) cross-section; (b,d) longitudinal section.
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Figure 4. Metallography of the Ti-6Al-4V treated by HIP: (a) T1P1, (b) T1P2, (c) T1P3, (d) T2P1, (e) T2P2, (f) T2P3, (g) T3P1, (h) T3P2, and (i) T3P3.
Figure 4. Metallography of the Ti-6Al-4V treated by HIP: (a) T1P1, (b) T1P2, (c) T1P3, (d) T2P1, (e) T2P2, (f) T2P3, (g) T3P1, (h) T3P2, and (i) T3P3.
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Figure 5. Statistics of α-phase lath width.
Figure 5. Statistics of α-phase lath width.
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Figure 6. SEM of the Ti-6Al-4V treated by HIP: (a) T1P1, (b) T1P2, (c) T1P3, (d) T2P1, (e) T2P2, (f) T2P3, (g) T3P1, (h) T3P2, and (i) T3P3.
Figure 6. SEM of the Ti-6Al-4V treated by HIP: (a) T1P1, (b) T1P2, (c) T1P3, (d) T2P1, (e) T2P2, (f) T2P3, (g) T3P1, (h) T3P2, and (i) T3P3.
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Figure 7. Statistics of β-phase volume fractions.
Figure 7. Statistics of β-phase volume fractions.
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Figure 8. EDX analysis of Ti-6Al-4V samples after HIP treatment at T2P3: (a) SEM image; (b) Ti content; (c) Al content; (d) V content.
Figure 8. EDX analysis of Ti-6Al-4V samples after HIP treatment at T2P3: (a) SEM image; (b) Ti content; (c) Al content; (d) V content.
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Figure 9. Density test results of specimens.
Figure 9. Density test results of specimens.
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Figure 10. Tensile properties statistics of SLM-formed Ti-6Al-4V after different HIP treatments.
Figure 10. Tensile properties statistics of SLM-formed Ti-6Al-4V after different HIP treatments.
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Figure 11. Tensile fracture surface of SLM-formed and HIP-treated specimens: (a) SLM; (b) T1P1; (c) T1P2; (d) T1P3; (e) T2P2; (f) T3P3.
Figure 11. Tensile fracture surface of SLM-formed and HIP-treated specimens: (a) SLM; (b) T1P1; (c) T1P2; (d) T1P3; (e) T2P2; (f) T3P3.
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Table
Table 1. Chemical compositions of Ti-6Al-4V titanium alloy powder (ω/%).
Table 1. Chemical compositions of Ti-6Al-4V titanium alloy powder (ω/%).
ElementAlVFeCHONTi
Content6.183.820.160.00750.00270.0690.012Bal.
Table
Table 2. The physical property of Ti-6Al-4V titanium alloy powder.
Table 2. The physical property of Ti-6Al-4V titanium alloy powder.
Tap DensityApp. DensityD(10) 1D(50) 2D(90) 3
2.85 g/cm32.44 g/cm326.043 μm39.405 μm58.975 μm
1 D(10) indicates the particle size corresponding to a cumulative particle size distribution of 10%. 2 D(50) indicates the particle size corresponding to a cumulative particle size distribution of 50%. 3 D(90) indicates the particle size corresponding to a cumulative particle size distribution of 90%.
Table
Table 3. Forming parameter of Ti-6Al-4V titanium alloy formed by SLM.
Table 3. Forming parameter of Ti-6Al-4V titanium alloy formed by SLM.
Laser PowerScanning SpeedHatching SpacingLayer Thickness
350 W1000 mm/s110 μm60 μm
Table
Table 4. Hot isostatic pressing schedules.
Table 4. Hot isostatic pressing schedules.
Sample LabelTemperaturePressure
SLM-formed//
T1P1910 °C110 MPa
T2P1930 °C110 MPa
T3P1950 °C110 MPa
T1P2910 °C130 MPa
T2P2930 °C130 MPa
T3P2950 °C130 MPa
T1P3910 °C150 MPa
T2P3930 °C150 MPa
T3P3950 °C150 MPa
T1, T2, and T3 stand for 910 °C, 930 °C, and 950 °C, respectively. P1, P2, and P3 stand for 110 MPa, 150 MPa, and 150 MPa, respectively.
Table
Table 5. Summary of tensile properties of SLM-formed Ti-6Al-4V after different HIP treatments.
Table 5. Summary of tensile properties of SLM-formed Ti-6Al-4V after different HIP treatments.
Sample Labelλα (μm)Vβ (%)YS (MPa)UTS (MPa)EL. (%)
SLM-formed//1103 ± 41226 ± 210.3 ± 0.6
T1P11.01813.07930 ± 11000 ± 116.5 ± 0.5
T2P11.11814.17909 ± 3982 ± 315.8 ± 0.8
T3P11.18420.58912 ± 2962 ± 216.5 ± 0.2
T1P20.91814.37935 ± 31005 ± 116.5 ± 0.5
T2P20.95615.07909 ± 5981 ± 415 ± 0.9
T3P21.14822.09891 ± 2967 ± 215.7 ± 0.8
T1P31.10215.96923 ± 5991 ± 415.3 ± 0.6
T2P31.13816.33908 ± 2976 ± 114.3 ± 0.8
T3P31.49227.85883 ± 5955 ± 315.2 ± 1
ASTM F2924-14//82589510
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MDPI and ACS Style

Lv, Z.; Li, H.; Che, L.; Chen, S.; Zhang, P.; He, J.; Wu, Z.; Niu, S.; Li, X. Effects of HIP Process Parameters on Microstructure and Mechanical Properties of Ti-6Al-4V Fabricated by SLM. Metals 2023, 13, 991. https://doi.org/10.3390/met13050991

AMA Style

Lv Z, Li H, Che L, Chen S, Zhang P, He J, Wu Z, Niu S, Li X. Effects of HIP Process Parameters on Microstructure and Mechanical Properties of Ti-6Al-4V Fabricated by SLM. Metals. 2023; 13(5):991. https://doi.org/10.3390/met13050991

Chicago/Turabian Style

Lv, Zhoujin, Haofeng Li, Lida Che, Shuo Chen, Pengjie Zhang, Jing He, Zhanfang Wu, Shanting Niu, and Xiangyang Li. 2023. "Effects of HIP Process Parameters on Microstructure and Mechanical Properties of Ti-6Al-4V Fabricated by SLM" Metals 13, no. 5: 991. https://doi.org/10.3390/met13050991

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