Microstructure and tensile fracture behavior of three-stage heat treated inconel 718 alloy produced via laser powder bed fusion process

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Original Article

Microstructure and tensile fracture behavior of three-stage heat treated inconel 718 alloy produced via laser powder bed fusion process Jun-Ren Zhao, Fei-Yi Hung ∗ , Truan-Sheng Lui Department of Materials Science and Engineering, National Cheng Kung University, Tainan 701, Taiwan

a r t i c l e

i n f o

a b s t r a c t

Article history:

In this study, Laser Powder Bed Fusion (LPBF) Inconel 718 is subjected to various heat

Received 19 October 2019

treatments, namely double aging, solid solution + double aging, and homogenization +

Accepted 9 January 2020

solid-solution + double aging, to investigate the effect of heat treatment on room- and high-

Available online xxx

temperature tensile properties. The results show that all three heat treatments increase

Keywords:

increase in tensile test temperature, the stress-induced Portevin-Le Chatelier (PLC) effect

hardness and room-temperature tensile strength, but greatly reduce ductility. With an Inconel 718

can effectively prevent the occurrence of oxidation at grain boundaries and maintain a cer-

Laser powder bed fusion (LPBF)

tain ductility in the range of room temperature to 600 ◦ C. However, when the temperature

High-temperature tensile properties

is 650 ◦ C, the PLC effect disappears, resulting in a high-temperature embrittlement effect.

Embrittlement Oxidation

1.

© 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http://creativecommons.org/licenses/by-nc-nd/4.0/).

Introduction

Inconel 718 (IN718) is a nickel-based superalloy that is strengthened by precipitating D022-ordered ␥” (Ni3 Nb) and L12-ordered ␥’ (Ni3 (Al, Ti)) phases in the ␥ matrix. Due to its excellent corrosion resistance and mechanical properties at room and elevated temperatures, IN718 is widely applied in energy, astronautic, and aeronautic industries for gas turbines, turbine blades, rocket engines, and nuclear power plant components [1–4]. However, its high hardness and low thermal conduction make cutting difficult [5].

∗

To reduce costs and mold consumption and create complex geometric shapes. Additive Manufacturing (AM) technology has been applied to various industries [6]. AM technologies are divided into two broad categories: (1) powder/wire feed systems [7,8], such as Directed Energy Deposition (DED) [9,10]; (2) powder bed systems, such as Electron Beam Melting (EBM) [11] and Laser Powder Bed Fusion (LPBF) [12,13,20–25]. In LPBF process, the metal powders are melted in a specified area with a high-energy laser beam, then it is rapidly solidified at a large cooling rate to build the component layer-by-layer [12,13]. These days, LPBF process is one of the major AM technologies for the production of nickel-based superalloys [14–19]. Many studies have investigated the effects of process parameters on the density and microstructure of as-prepared LPBF IN718 [20,21]. The effects of heat treatment on the microstructure and mechanical properties of LPBF IN718 have also been investigated [22–25]. The differences between the

Correspondence author. E-mail: [email protected] (F. Hung). https://doi.org/10.1016/j.jmrt.2020.01.030 2238-7854/© 2020 The Authors. Published by Elsevier B.V. This is an open access article under the CC BY-NC-ND license (http:// creativecommons.org/licenses/by-nc-nd/4.0/). Please cite this article in press as: Zhao J, et al. Microstructure and tensile fracture behavior of three-stage heat treated inconel 718 alloy produced via laser powder bed fusion process. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.030

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Fig. 2 – Dimensions of a tensile specimen.

Fig. 1 – Macroscopic morphology photographs of as-printed IN718 specimens.

LPBF process and traditional forging and casting of IN718 under standard heat treatment conditions have been discussed [25]. The hardness and strength of LPBF IN718 increase after general industrial heat treatment [26,27]. Although IN718 is widely used in high- temperature environments, few studies have evaluated the high-temperature mechanical properties of LPBF IN718 [28–30]. It has been found that LPBF IN718 gradually decreases in strength as the tensile test temperature rises, and that ductility depends on the direction of printing [28,29]. The cause of the deterioration of high-temperature elongation has not been discussed. According to previous research [25–27], the strength of LPBF IN718 is insufficient for direct industrial applications. It is hoped that strength and ductility can be improved by appropriate heat treatment to meet the requirements of room- and high-temperature applications. In the present study, three heat treatment parameters were selected to investigate the high-temperature tensile properties of LPBF IN718. High-temperature aging and the high-temperature failure mechanism are evaluated to clarify the embrittlement effect of high temperature on LPBF IN718. In general, metals exhibit decreased strength and increased ductility at high-temperature; in contrast, LPBF IN718 exhibits decreased strength and ductility. This study clarifies the hightemperature embrittlement mechanism of LPBF IN718. The results can be used as a reference for aeronautic and astronautic applications.

2.

Experimental procedure

IN718 powder particles were purchased from Electro-Optical Systems (EOS) GmbH (Krailling, Germany). Their chemical composition is Ni (50−55 wt.%), Cr (17−21 wt.%), Nb (4.7–5.5 wt.%), Mo (2.8–3.3 wt.%), Ti (0.6–1.2 wt.%), Al (0.2–0.8 wt.%), Co (≤1.0 wt.%), Cu (≤0.3 wt.%), C (≤0.08 wt.%), and Fe (bal.). The LPBF process parameters used in this study are shown in Table 1. The specimens were fabricated using an EOS M290 400 W machine (EOS, Krailling, Germany) in an inert gas (argon) atmosphere, as shown in Fig. 1, and then removed from their supports using electrical discharge machining wire cutting.

The as-prepared LPBF IN718 specimens, denoted AS, were subjected to various different heat treatments, namely double aging heat treatment (720 ◦ C, 8 h/furnace cooling at 55 ◦ C/h to 620 ◦ C, 8 h/air cooling), solid solution (980 ◦ C, 1 h/water cooling) + double aging heat treatment (720 ◦ C, 8 h/furnace cooling at 55 ◦ C/h to 620 ◦ C, 8 h/air cooling), and homogenization (1080 ◦ C, 1.5 h/water cooling) + solid solution (980 ◦ C, 1 h/water cooling) + double aging heat treatment (720 ◦ C, 8 h/furnace cooling at 55 ◦ C/h to 620 ◦ C, 8 h/air cooling). The specimens obtained using these treatments are denoted as A, SA, and HSA, respectively. The specimens were polished with SiC paper (from #120 to #5000), 1- and 0.3-␮m Al2 O3 aqueous solution, and 0.04-␮m SiO2 polishing solution, and then etched with a chemical solution consisting of 50% HCl, 10% HNO3 , 2% HF, and 38% distilled water to examine the microstructure. Optical Microscopy (OM, OLYMPUS BX41M-LED, Tokyo, Japan) and Scanning Electron Microscopy (SEM, HITACHI SU-5000, HITACHI, Tokyo, Japan) were used to examine the microstructure, and X-ray diffraction (XRD, Bruker AXS GmbH, Karlsruhe, Germany) was used to analyze the phase composition. The dimensions of a tensile specimen are shown in Fig. 2. A universal testing machine (HT-8336, Hung Ta, Taichung, Taiwan) was used to perform the tensile tests. The crosshead speed was 1 mm/min, corresponding to an initial strain rate of 8.33 × 10−4 s−1 . The AS, A, SA, and HSA specimens were subjected to room- and high-temperature (650 ◦ C) tensile tests to analyze the mechanical properties of LPBF IN718. In addition to the observation of the tensile fracture characteristics, the SA specimen with the best mechanical properties was selected to carry out tensile tests at various temperatures (500, 550, and 600 ◦ C) to confirm the critical temperature of hightemperature embrittlement. At least three specimens were used for each test. The mean value for the test specimens was taken as the result for the corresponding condition. Finally, the SA specimen was polished using # 80 to #1000 SiC paper to remove the surface oxide layer and subjected to high-temperature oxidation at 650 ◦ C for 0.5 h to simulate the environment in which the high-temperature tensile test piece was placed. The phase structure and aging characteristics of the as-prepared SA specimen, surface-layer-polished specimen, and high-temperature oxidation specimen were analyzed using XRD (Bruker AXS GmbH, Karlsruhe, Germany). Furthermore, the O, Cr, and Ni element concentration distribution along the longitudinal directions in the high-temperature oxidation specimen at 600 ◦ C and 650 ◦ C for 0.5 h were

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Table 1 – Process parameters used for the Laser Powder Bed Fusion process. Laser power 230 W

Scanning velocity 760 mm/s

Layer thickness 40 ␮m

Hatching space 110 ␮m

Idle time 10 s

Preheat temperature ◦

80 C

Porosity 99.7 %

Fig. 3 – (a) 3D metallographic microstructure of the as-prepared LPBF IN718 specimen, (b) low-magnification SEM morphology in the side direction, and (c) magnified SEM morphology in the side direction.

measured using Secondary Ion Mass Spectrometry (SIMS) to evaluate the oxide layer characteristics.

3.

Results and discussion

3.1.

Microstructure evolution

The metallographic microstructure of the as-prepared LPBF IN718 specimen viewed in three-dimensional (3D) sections is shown in Fig. 3(a). It can be seen that the Side Direction (SD) plane (i.e., direction vertical to the laser direction) shows melt pool stacked morphologies. Some long strip morphologies (laser scanning direction) can be observed in the Normal Direction (ND) plane (i.e., direction parallel to the laser direction). Fig. 3(b) shows the microstructure characteristics of the SD plane. Columnar dendrites grow nearly parallel to the laser direction with a dendrite arm spacing of about 200–600 nm. In addition, the columnar dendrites grow in different directions in a single melt pool. The matrix is gray and the dendritic grain boundary of solidification segregation is white (Fig. 3c). Fig. 4 shows the SD microstructure of the LPBF IN718 subjected to various heat treatments. Fig. 4(a) shows that the as-prepared LPBF specimen has a layered stacked melt pool structure with a diameter of approximately 150 ␮m. After double aging heat treatment (A specimen), the melt pool traces become inconspicuous (effect of thermal diffusion), and long equiaxed grains penetrate the melt pools, as shown in Fig. 4(b).

The microstructure of the specimen subjected to solid solution treatment + double aging heat treatment (SA specimen) is shown in Fig. 4(c). The melt pool traces and the dendritic structure completely disappear, and are replaced by long equiaxed grains in the matrix parallel to the laser direction. This indicates that after the solid solution heat treatment, the decomposed dendrites and the solute atoms in the segregation region are re-solidified back into the matrix and form the recrystallized long equiaxed grains. Fig. 4(d) shows the microstructure of the specimen subjected to homogenization treatment + solid solution treatment + double aging heat treatment (HSA specimen). The HSA specimen has a completely equiaxed grain matrix; the grains are more circularly passivated than those in Fig. 4(c). Previous studies found similar IN718 microstructure after homogenization treatment [30–32]. The microstructures of the specimens subjected to three heat treatment processes in this study are significantly different (AS specimen: columnar-dendrites + melt pool phase boundaries; A specimen: dendrites + residual part of melt pool phase boundaries + long equiaxed grains; SA specimen: long equiaxed grains; HSA specimen: circularly passivated equiaxed grains). Fig. 5 shows the high magnification SEM image of HSA specimen. Spherical and plate ␥’ particle precipitates can be observed within the matrix (40−80 nm). The ␥” precipitates are too small, so they cannot be seen. Fig. 6 shows the XRD patterns of the LPBF IN718 specimens subjected to various heat treatments. The as-prepared LPBF IN718 mainly has (111), (200), (220), and (311) peaks, the peaks

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Fig. 4 – Microstructures in the side direction of (a) AS, (b) A, (c) SA, and (d) HSA specimens.

Fig. 5 – Magnified SEM morphology of HSA specimen matrix.

corresponding to ␥, ␥’, and ␥” almost overlap. According to previous reports [30,33], the (111) and (222) peaks correspond to the ␥ and ␥’ phases, respectively. Therefore, the as-prepared LPBF IN718 contained the precipitation-strengthening phases ␥’ and ␥” in addition to the ␥ matrix. The XRD patterns of A specimen were similar to as-prepared LPBF IN718. Compared with the as-prepared LPBF specimen, the SA specimen (solid solution and double aging) had the ␦ phase, Cr2 O3 , and MC carbides. Compared with the SA specimen, the content levels

Fig. 6 – XRD patterns of AS, A, SA, and HSA specimens.

of the ␦ phase, Cr2 O3 , and MC carbides of the HSA specimen (homogeneous, solid solution, and double aging heat treatment) were higher. This indicates that heat treatment at above 980 ◦ C causes oxidation of the material and the formation of MC carbides.

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Fig. 7 – Hardness of AS, A, SA, and HSA specimens.

3.2.

Room-temperature tensile properties

Fig. 7 shows the hardness comparison of the AS, A, SA, and HSA specimens. Due to the high cooling rate in the LPBF process, there is insufficient time for the complete precipitation of the strengthening phases [25], resulting in lower hardness. After double aging heat treatment, the strengthening phases precipitate, greatly improving hardness. The hardness of the A specimen was similar to those of the SA and HSA specimens, which indicates that the double aging heat treatment had a dominant effect on hardness. The room-temperature tensile properties of the AS, A, SA, and HSA specimens are shown in Fig. 8. The AS specimen had excellent ductility (total elongation: nearly 30%), but its strength was lower than the ASM 5662 standard [34]. Furthermore, the three heat treatments significantly improved strength, indicating that the double aging heat treatment can effectively improve tensile strength. The strengths of the SA and HSA specimens were slightly lower than that of the A specimen (only aging heat treatment). The ductility of the heat-treated specimens decreased significantly. It is worth noting that the ductility of the SA and HSA specimens are better than that of the A specimen and higher than ASM 5662. The ductility of the SA specimen is the best among the three heat-treated specimens; the A specimen exhibited a significant room-temperature embrittlement effect. Fig. 9 shows macroscopic morphology photographs of room temperature tensile fracture specimens (from left to right: AS, A, SA, and HSA specimens). The AS specimen, which had the best ductility, shows obvious shrinkage and elongation; the three heat-treated specimens show similar behavior. It is worth noting that the surface color is different for the four specimens, which indicates different surface films. To further investigate the effect of heat treatment on the tensile mechanical properties, the morphology of the tensile fracture sub-surfaces after heat treatment is shown in Fig. 10. The fracture sub-surface of the AS specimen, shown in Fig. 10(a), shows melt pool ductile fracture morphologies, with fine recrystallized structures near the fracture. The fracture subsurface of the A specimen is flatter (cleavage feature) and the crack is directly through the grains (i.e., a transgranular

Fig. 8 – Tensile properties of AS, A, SA, and HSA specimens at room temperature. (a) Strength and (b) ductility. (YS: yield strength; UTS: ultimate tensile strength; UE: uniform elongation; TE: total elongation.).

Fig. 9 – Macroscopic morphology photographs of room-temperature tensile fracture specimens. fracture) (Fig. 10b). The main reason for this crack is that the matrix composition is complex (dendrites + residual part of melt pool phase boundaries + long equiaxed grains), resulting in the embrittlement effect. The fracture subsurface of the SA specimen shown in Fig. 10(c) is similar to that in Fig. 10(d), but

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Fig. 10 – Morphologies of tensile fracture subsurfaces at room temperature for (a) AS, (b) A, (c) SA, and (d) HSA specimens.

many small and mainly long equiaxed grains appear improving ductility. The ductility of the HSA specimen is slightly lower than that of the SA specimen because of bigger circularly passivated equiaxed grains (Fig. 10d). The AS specimen has excellent ductility (>25%), but its strength is lower than the general industry application standards [23,25,29]. The three heat treatments, increased the strength to above the application standards. Because the A specimen was directly subjected to by double aging heat treatment, part of the strengthening phase precipitated between dendritic arms, resulting in a significant decrease in ductility [25,35]. For the SA specimen, which was subjected to solid solution treatment, the solid solution atoms first dissolved into the matrix and then precipitated. The matrix had long equiaxed grains and the strengthening phases were evenly distributed, making the ductility of the SA specimen better than that of the A specimen. The HSA specimen was subjected to homogenization treatment, which led to equiaxed grain growth. Due to the contribution of the precipitated phase and the oxidized phase, the ductility was slightly lower than that of the SA specimen. For room-temperature applications, the SA and HSA specimens meet the requirements.

3.3.

High-temperature embrittlement

High-temperature strength is an important indicator of 718 nickel-based superalloys. In this study, the high-temperature

(650 ◦ C; industrial application >648 ◦ C) tensile properties of the specimens are investigated to evaluate the hightemperature embrittlement mechanism. Fig. 11 shows the high-temperature tensile properties of the AS, A, SA, and HSA specimens. Compared with Fig. 8, the strengths of all specimens are lower at 650 ◦ C. The ductility (Fig. 11b) of all specimens except for the AS specimen, which increased, significantly decreased, indicating high-temperature embrittlement. Fig. 12 shows macroscopic morphology photographs of high temperature (650 ◦ C) tensile fracture specimens (from left to right: AS, A, SA, and HSA specimens). The AS specimen shows obvious shrinkage and elongation; the three heat-treated specimens show no obvious shrinkage and their ductility is similar. It is worth noting that the SA specimen had better elongation and that its fracture was at 45◦ . To understand the cause of high-temperature embrittlement, the fracture subsurface of the high-temperature tensile specimens is investigated. The high-temperature fracture subsurface of the AS specimen is similar to that of the room temperature subsurface, showing a ductile fracture morphology (Fig. 13a). Fig. 13(b) shows the fracture subsurface of the A specimen, which has a zigzag shape; there are small melt pools or dendrites near the fracture. The SA specimen has a zigzag pattern fracture and an increasing grain size, as shown in Figs. 13(c) and 13(d). The HSA specimen also has a zigzag pattern fracture (Fig. 13e).

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Fig. 12 – Macroscopic morphology photographs of high-temperature (650 ◦ C) tensile fracture specimens.

lier (PLC) effect. It is worth noting that as the temperature rises, the serrations become larger and then completely disappear at 650 ◦ C. This temperature marks the end of the PLC effect, and is thus the critical temperature of the high-temperature embrittlement of the SA specimen.

3.4.

Fig. 11 – Tensile properties of AS, A, SA, and HSA specimens at high temperature (650 ◦ C). (a) Strength and (b) ductility.

The crack propagation is vertical to the tensile direction, as shown in Fig. 13(f). This indicates that the grain boundary of the equiaxed grains is the crack initiation point at high temperature, and that crack growth is scattered. Furthermore, during the tensile test at 650 ◦ C, grain growth and precipitation behavior were found for each specimen; the fracture characteristics changed from a cleavage feature to a zigzag pattern fracture. In general, at high temperatures, the strength of a metal decreases and ductility increases. In this study, LPBF IN718 showed a high-temperature embrittlement effect. To investigate the cause of the high-temperature embrittlement and the corresponding temperature range, the SA specimen, which had the best high-temperature ductility, was selected to investigate properties at various temperatures. Fig. 14 shows the different temperature tensile properties of the SA specimen. As the tensile test temperature increases, the ductility remains at a certain value up to 600 ◦ C, but then drops significantly at 650 ◦ C. This means that 650 ◦ C is the critical temperature for the high-temperature embrittlement of IN718. According to the high-temperature tensile stress-strain curve (Fig. 15), serrated flow is observed in the temperature range of 500−600 ◦ C. According to previous research [35,36], this serrated flow can be attributed to the Portevin–Le Chate-

High-temperature oxidation and over aging

Fig. 16 shows the XRD patterns of the as-prepared SA specimen, the surface-layer-polished SA specimen, and the SA specimen subjected to high-temperature oxidation at 650 ◦ C for 0.5 h in an air furnace. The Cr2 O3 and MC carbide peaks of the surface-layer-polished SA specimen disappear, indicating that the high-temperature heat treatment led to the formation of oxides and carbides on the surface. For the SA specimen subjected to high-temperature oxidation at 650 ◦ C for 0.5 h, a weak Cr2 O3 peak appears. This indicates that oxidation occurs at 650 ◦ C, but carbides do not form (they need a long time to precipitate). To confirm the oxidation degree of the SA specimen at 650 ◦ C and the causes of the high-temperature ductility difference at 600 ◦ C and 650 ◦ C, SIMS analysis (Fig. 17) is performed on the specimens for these two temperatures. The thickness of the oxide layer at 650 ◦ C was thicker than that at 600 ◦ C, and the oxygen signal was not only on the surface but also inside the material. This result is consistent with the XRD patterns (Fig. 16). According to previous research [31–40], for IN718 in a high-temperature oxygen environment, oxygen easily diffuses into the grain boundary, decreasing ductility (grain boundary embrittlement). According to one study [41], the free energy of Cr2 O3 is low and thus Cr2 O3 easily reacts at the grain boundary of the surface layer. Therefore, the specimen easily cracks along the grains during high-temperature deformation. The PLC effect that causes the serrated flow of the stress-strain curve can inhibit oxidation at the grain boundary, decreasing high-temperature embrittlement [31,40]. Furthermore, this study found that 650 ◦ C is the PLC effect termination temperature, and is thus the critical temperature

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Fig. 13 – Morphologies of tensile fracture sub-surfaces at room temperature for (a) AS, (b) A, (c) SA, (d) crack of SA, (e) HSA, and (f) crack of HSA specimens.

of high-temperature embrittlement. Considering that 650 ◦ C is higher than the double aging temperature of 620 ◦ C, and that grain growth occurs at a high-temperature tensile fracture structure, it can be confirmed that the overaging occurs at 650 ◦ C.Fig. 18 shows a schematic diagram of the high-temperature embrittlement mechanism for LPBF IN718. When the temperature is higher than 650 ◦ C, the material exhibits grain boundary oxidation and overaging, with no PLC effect. This decreases tensile ductility and the high-temperature grain boundary embrittlement effect.

4.

Conclusions

1 The AS specimen shows layered melt pool stacked morphologies, columnar dendrites grow nearly parallel to the laser direction. After heat treatments, the A specimen shows dendrites and long equiaxed grains with residual part of melt pool phase boundaries; SA specimen and HSA specimen show the long and circularly passivated recrystallized equiaxed grains, respectively. In addition,

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Fig. 16 – XRD patterns of as-prepared SA, polished SA, and SA subjected to 650 ◦ C for 0.5 h specimens.

Fig. 14 – High-temperature tensile properties of SA specimen. (a) Strength and (b) ductility.

Fig. 17 – SIMS results of high-temperature oxidization at 600 ◦ C and 650 ◦ C for 0.5 h of SA specimen.

Fig. 15 – High-temperature stress-strain curve of SA specimen.

Cr2 O3 film is formed on the surface of the SA and HSA specimens. 2 Heat treatment processes (A, SA, and HSA) increase the tensile strength and decrease the ductility of LPBF IN718 at room-temperature. The double aging treatment process improved the hardness and tensile strength. The A

specimen exhibited significant room-temperature embrittlement. The SA and HSA specimens had recrystallized equiaxed structures to maintain a certain value of ductility. 3 650 ◦ C is the critical temperature of the high-temperature embrittlement of LPBF IN718. Below this temperature, PLC effect can effectively prevent the occurrence of oxidation at grain boundaries and maintain a certain ductility. Above this temperature, grain boundary oxidation occurs and overaging behavior is induced, resulting in intergranular cracking and high-temperature brittleness, and thus a deterioration in high-temperature tensile ductility.

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[3]

[4]

[5]

[6]

Fig. 18 – Schematic diagram of high-temperature embrittlement mechanism.

Conflict of interest We declare that we have no financial and personal relationships with other people or organizations that can inappropriately influence our work, there is no professional or other personal interest of any nature or kind in any product, service and/or company that could be construed as influencing the position presented in, or the review of, the manuscript entitled, “Microstructure and Tensile Fracture Behavior of Three-Stage Heat Treated Inconel 718 Alloy Produced via Laser Powder Bed Fusion Process”.

[7]

[8]

[9]

[10]

Acknowledgments [11]

The authors are grateful to the Instrument Center of National Cheng Kung University and the Ministry of Science and Technology of Taiwan (Grant No. MOST 1072221-E-006-012-MY2) for their financial support of this research.

[12]

Appendix A. Supplementary data Supplementary material related to this article can be found, in the online version, at doi:https://doi.org/10.1016/j. jmrt.2020.01.030.

[13]

[14]

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Please cite this article in press as: Zhao J, et al. Microstructure and tensile fracture behavior of three-stage heat treated inconel 718 alloy produced via laser powder bed fusion process. J Mater Res Technol. 2020. https://doi.org/10.1016/j.jmrt.2020.01.030