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metals Article Normalizing Effect of Heat Treatment Processing on 17-4 PH Stainless Steel Manufactured by Powder Bed Fusion 1 , 2 1 1 3 2 1 Si-Mo Yeon , Jongcheon Yoon , Tae Bum Kim , Seung Ho Lee , Tea-Sung Jun , Yong Son 4 , and Kyunsuk Choi * Advanced Joining & Additive Manufacturing R&D Department, Korea Institute of Industrial Technology, Siheung 15014, Korea; simo@kitech.re.kr (S.-M.Y.); jongcheon89@kitech.re.kr (J.Y.); xoqja78@kitech.re.kr (T.B.K.); sonyong@kitech.re.kr (Y.S.) Department of Mechanical Engineering, Incheon National University, Yeonsu, Incheon 22012, Korea; t.jun@inu.ac.kr HYU-KITECH Joint Department, Hanyang University, Ansan 15588, Korea; qhtmho@hanyang.ac.kr Department of Industry-University Convergence, Hanbat National University, Yuseong, Daejeon 34158, Korea * Correspondence: kschoi@hanbat.ac.kr; Tel.: +82-42-828-8586; Fax: +82-42-821-1585 Abstract: Laser powder bed fusion (L-PBF)-processed 17-4 PH stainless steel (SS) generally exhibits a non-equilibrium microstructure consisting mostly of columnar -ferrite grains and a substantial fraction of retained austenite and martensite, contrary to 17-4 PH SS wrought with a fully martensite structure and coarse grains. Despite the different microstructures of L-PBF and wrought 17-4 PH SS, post-processing is typically performed using the conventional heat treatment method. The insufficient effect of the heat treatment on the L-PBF product produces a -ferrite phase in the microstructure. To obtain improved mechanical properties, the addition of a normalizing treatment to the conventional heat treatment after L-PBF in a nitrogen gas environment was investigated. The fully martensitic matrix developed by adding the normalizing treatment contained homogeneous Cu precipitates and exhibited a similar or improved strength and elongation to failure compared to the wrought SS. Citation: Yeon, S.-M.; Yoon, J.; Kim, T.B.; Lee, S.H.; Jun, T.-S.; Son, Y.; Choi, Keywords: 17-4 PH stainless steel; selective laser melting; laser powder bed fusion; heat treatment; K. Normalizing Effect of Heat microstructure; mechanical properties Treatment Processing on 17-4 PH Stainless Steel Manufactured by Powder Bed Fusion. Metals 2022, 12, 704. https://doi.org/10.3390/ met12050704 1. Introduction A 17-4 PH stainless steel (SS) is a chromium-nickel-copper precipitation-hardenable Academic Editor: Marco Mandolini alloy that is widely used in the aerospace, marine, power generation, chemical plant, and Received: 24 March 2022 medical industries. It requires high strength and corrosion resistance against thermal expo- Accepted: 16 April 2022 sure to temperatures of approximately 316 C [1,2]. When fabricated using the conventional Published: 20 April 2022 process, this alloy is predominantly composed of a martensite phase and is strengthened Publisher’s Note: MDPI stays neutral by heat treatment (aging treatment at 482 C for 1 h after solution treatment at 1060 C) with regard to jurisdictional claims in to precipitate Cu particles in the martensite [3,4]. Recently, higher value-added indus- published maps and institutional affil- tries, such as aerospace, health care, power generation, and automobiles, have demanded iations. more design freedom and complex manufacturing technology to improve functions and performance [5,6]. Additive manufacturing technology has attracted attention as a global solution, with the functional enhancement of dissimilar materials, such as superior corrosion resistance, Copyright: © 2022 by the authors. excellent thermal stability, high hardness, better wear performance, [7–9], and the fabrica- Licensee MDPI, Basel, Switzerland. tion of near-net-shape parts with higher dimensional accuracy [10]. This article is an open access article A 17-4 PH SS is comparatively difficult to work with and has poor machinability, distributed under the terms and thereby leading to a demand for additive manufacturing technology for the fabrication of conditions of the Creative Commons 17-4 PH SS products, with customized shapes for specialized applications [11]. Previous Attribution (CC BY) license (https:// studies have shown that the characteristics of laser powder bed fusion (L-PBF) products are creativecommons.org/licenses/by/ strongly dependent on the materials (powder composition, initial atomizing media, etc.), 4.0/). Metals 2022, 12, 704. https://doi.org/10.3390/met12050704 https://www.mdpi.com/journal/metals Metals 2022, 12, x FOR PEER REVIEW 2 of 14 Metals 2022, 12, 704 2 of 14 17-4 PH SS products, with customized shapes for specialized applications [11]. Previous studies have shown that the characteristics of laser powder bed fusion (L-PBF) products are strongly dependent on the materials (powder composition, initial atomizing media, the build-chamber atmosphere, the design parameters, the process conditions, and the etc.), the build-chamber atmosphere, the design parameters, the process conditions, and post-heat treatments [12]. the post-heat treatments [12]. Additively manufactured 17-4 PH SS has a unique microstructure that affects its me- Additively manufactured 17-4 PH SS has a unique microstructure that affects its me- chanical properties and phase evaluations during post-processing. It exhibits an anisotropic chanical properties and phase evaluations during post-processing. It exhibits an aniso- microstructure with columnar grains that are oriented parallel to the build direction and tropic microstructure with columnar grains that are oriented parallel to the build direction 5 7 fine grains with average sizes smaller than 10 m when cooled quickly at 10 –10 K/s [13]. 5 7 and fine grains with average sizes smaller than 10 μm when cooled quickly at 10 –10 K/s It also contains non-equilibrium microstructures comprising mixed phases of -ferrite [13]. It also contains non-equilibrium microstructures comprising mixed phases of δ-fer- (BCC), austenite (FCC, ), martensite (BCT, ’), and carbide precipitations [11]. rite (BCC), austenite (FCC, γ), martensite (BCT, α΄), and carbide precipitations [11]. The presence of retained austenite and -ferrite phases and copper precipitates in the The presence of retained austenite and δ-ferrite phases and copper precipitates in the martensite matrix and the changes in grain size that occur during the manufacturing process martensite matrix and the changes in grain size that occur during the manufacturing pro- significantly influence the mechanical properties of the fabricated parts. The austenite cess significantly influence the mechanical properties of the fabricated parts. The austenite phase reduces the amount of copper precipitated during the aging step because of the phase reduces the amount of copper precipitated during the aging step because of the higher solubility and lower diffusivity of copper in austenite than in martensite [14–16]. higher solubility and lower diffusivity of copper in austenite than in martensite [14–16]. The remaining ferrite also exhibits delayed hardening behavior [17,18]. The fraction of The remaining ferrite also exhibits delayed hardening behavior [17,18]. The fraction of austenite and ferrite remaining after the L-PBF process and heat treatment depends on the austenite and ferrite remaining after the L-PBF process and heat treatment depends on the build-chamber environment (N or Ar gas) [19], grain size (which may be reduced by using build-chamber environment (N2 or Ar gas) [19], grain size (which may be reduced by us- a high cooling rate) [20,21], residual stress in the grains or at the grain boundaries [19,22,23], ing a high cooling rate) [20,21], residual stress in the grains or at the grain boundaries composition (chromium-to-nickel equivalent value) [24], and the presence of austenite [19,22,23], composition (chromium-to-nickel equivalent value) [24], and the presence of stabilizers such as C and N [25]. austenite stabilizers such as C and N [25]. Figure 1 shows the mechanical properties of the as-built (AB) specimens produced by Figure 1 shows the mechanical properties of the as-built (AB) specimens produced the L-PBF and heat-treatment (HT) specimens collected from the literature [23,26–30] and by the L-PBF and heat-treatment (HT) specimens collected from the literature [23,26–30]. the detailed data of mechanical properties in Figure 1 is presented in the Supplementary The graph shows that the yield strength of the AB L-PBF parts is less than 800 MPa, and Table S1. The graph shows that the yield strength of the AB L-PBF parts is less than 800 MPa, the yield strength and elongation of the HT parts are less than the corresponding values and the yield strength and elongation of the HT parts are less than the corresponding values of the wrought parts, for which ASTM A693 (H900 condition) is the standard reference of the wrought parts, for which ASTM A693 (H900 condition) is the standard reference [31]. [31]. The reduced yield strength of the L-PBF 17-4 PH SS products is important because it The reduced yield strength of the L-PBF 17-4 PH SS products is important because it is used is used to determine the safety coefficient in most structural designs. The mechanical to determine the safety coefficient in most structural designs. The mechanical properties of properties of the L-PBF products should be improved, and further research on additive the L-PBF products should be improved, and further research on additive processing and processing and post-process optimization is required. post-process optimization is required. Figure Figure 1. 1. Mech Mechanical anical properties of 17-4 PH SS produc properties of 17-4 PH SS produced ed by by L-PBF from previous studies [ L-PBF from previous studies [23 23,26–31]. ,26–31]. Further research is required to extend the industrial application of the L-PBF process because it is not sufficiently economical in many cases. The manufacturing cost of the additive manufacturing process exceeds that of most traditional methods [32–36]. In Metals 2022, 12, x FOR PEER REVIEW 3 of 14 Further research is required to extend the industrial application of the L-PBF process Metals 2022, 12, 704 3 of 14 because it is not sufficiently economical in many cases. The manufacturing cost of the ad- ditive manufacturing process exceeds that of most traditional methods [32–36]. In previ- ous studies, the effects of the mechanical properties of Ar and N atmospheres during the previous studies, the effects of the mechanical properties of Ar and N atmospheres during L-PBF process were studied, and it is generally reported that poor mechanical properties the L-PBF process were studied, and it is generally reported that poor mechanical properties appear in a nitrogen atmosphere [12,15,19]. appear in a nitrogen atmosphere [12,15,19]. Motivated by this, we studied the microstructure and mechanical properties based Motivated by this, we studied the microstructure and mechanical properties based on the additional heat treatment and heat treatment steps of 17-4 PH SS fabricated under on the additional heat treatment and heat treatment steps of 17-4 PH SS fabricated under nitrogen conditions, which is advantageous in terms of cost, although it has unfavorable nitrogen conditions, which is advantageous in terms of cost, although it has unfavor- mechanical properties. Therefore, a heat treatment method for obtaining a stable micro- able mechanical properties. Therefore, a heat treatment method for obtaining a stable structure and predictable properties is proposed. microstructure and predictable properties is proposed. 2. Materials and Methods 2. Materials and Methods 2.1. 17-4 PH SS Powder 2.1. 17-4 PH SS Powder Gas-atomized 17-4 PH SS powders (obtained from 3D-systems Inc., Rock Hill, SC, Gas-atomized 17-4 PH SS powders (obtained from 3D-systems Inc., Rock Hill, SC, USA) with particle sizes ranging between 5–30 μm and with an average diameter of 10.63 USA) with particle sizes ranging between 5–30 m and with an average diameter of μm were used in this study, as shown in Figure 2. The chemical composition of the pow- 10.63 m were used in this study, as shown in Figure 2. The chemical composition of the ders powders is listed is listed in Table in T 1. able Inductively coupl 1. Inductively ed p coupled lasma-atom plasma-atomic ic emission spec emission trosco spectr py (ICP oscopy - AES (ICP-AES) ) and car and bon-su carbon-sulfur lfur analy analysis sis confconfirmed irmed that that the the 17-17-4 4 PH PH SS p SSo powder wder had hada anom nominal inal chemical composition that conformed to the 17-4 PH SS specification (ASTM A693-16). chemical composition that conformed to the 17-4 PH SS specification (ASTM A693-16). Figure 2. (a) SEM image and (b) particle size distribution of 17-4 PH stainless steel powder. Figure 2. (a) SEM image and (b) particle size distribution of 17-4 PH stainless steel powder. Table 1. Chemical compositions (wt.%) of 17-4 PH stainless steel powder. Table 1. Chemical compositions (wt.%) of 17-4 PH stainless steel powder. Fe Cr Ni Cu Si Mn Nb P S N C Fe Cr Ni Cu Si Mn Nb P S N C Bal. 17.10 4.40 4.42 0.92 0.85 0.24 0.021 0.009 0.009 0.027 Bal. 17.10 4.40 4.42 0.92 0.85 0.24 0.021 0.009 0.009 0.027 2.2. Part Fabrication by L-PBF 2.2. Part Fabrication by L-PBF The 17-4 PH SS test specimens were produced using a 3D-systems ProX300 machine. The 17-4 PH SS test specimens were produced using a 3D-systems ProX300 machine. The specimens were fabricated using a custom-built method, as shown in Figure 3a. The The specimens were fabricated using a custom-built method, as shown in Figure 3a. The laser scanning strategy for all the specimens consisted of an outer contour area and an laser scanning strategy for all the specimens consisted of an outer contour area and an inner hatching area. The contour was scanned in two lines at intervals of 40 μm, and the inner hatching area. The contour was scanned in two lines at intervals of 40 m, and the hatching in a 10 mm regular hexagonal pattern overlapped the remaining area by 50 μm, hatching in a 10 mm regular hexagonal pattern overlapped the remaining area by 50 m, as shown in Figure 3b. All specimens were fabricated with the same parameters: laser as shown in Figure 3b. All specimens were fabricated with the same parameters: laser power of 135 W, laser scan speed of 1200 mm/s, layer thickness of 40 μm, and hatching power of 135 W, laser scan speed of 1200 mm/s, layer thickness of 40 m, and hatching distance of 50 μm. The oxygen content in the build chamber was reduced to 1000 ppm to distance of 50 m. The oxygen content in the build chamber was reduced to 1000 ppm to prevent specimen contamination. Nitrogen gas was used as the inert instead of argon gas, prevent specimen contamination. Nitrogen gas was used as the inert instead of argon gas, for environmental control of the build chamber during the L-PBF process. for environmental control of the build chamber during the L-PBF process. Metals 2022, 12, x FOR PEER REVIEW 4 of 14 Metals 2022, 12, 704 4 of 14 Figure 3. Schematic of (a) L-PBF machine and (b) strategies of laser scanning and layering. Figure 3. Schematic of (a) L-PBF machine and (b) strategies of laser scanning and layering. 2.3. Post-Heat Treatment Procedures 2.3. Post-Heat Treatment Procedures The details of the post-heat-treatment procedures used in this study are listed in The details of the post-heat-treatment procedures used in this study are listed in Ta- Table 2. To analyze the effects of thermal exposure on microstructure and phase evolution, ble 2. To analyze the effects of thermal exposure on microstructure and phase evolution, the specimens were heat treated using conventional methods (ASTM A693, Solution A the specimens were heat treated using conventional methods (ASTM A693, Solution A + + H900, aging treatment at 482 C for 1 h, and after solution treatment at 1060 C) and H900, aging treatment at 482 °C for 1 h, and after solution treatment at 1060 °C) and sub- subjected to normalizing treatment (1200 C for 4 h before the solid solution treatment). jected to normalizing treatment (1200 °C for 4 h before the solid solution treatment). Five Five types of heat-treated specimens subjected to different numbers of treatment steps were types of heat-treated specimens subjected to different numbers of treatment steps were prepared. prepared. Table 2. Heat treatment condition of sample using fabricated L-PBF. Table 2. Heat treatment condition of sample using fabricated L-PBF. Sample Name Heat Treatment Condition Sample Name Heat Treatment Condition As-built (AB) - As-built (AB) - Normalizing treatment (N) 1200 C, 4 h + Furnace cooling Normalizing treatment (N) 1200 °C, 4 h + Furnace cooling Solution treatment (S) 1060 C, 1 h + Gas cooling Solution treatment (S) 1060 °C, 1 h + Gas cooling Solution + Aging treatment (SA) 1060 C, 1 h + Gas cooling ! 482 C, 4 h + Air cooling Solution + Aging treatment (SA) 1060 ℃, 1 h + Gas cooling → 482 °C, 4 h + Air cooling Normalizing + Solution + Aging 1200 C, 4 h + Furnace cooling ! 1060 C, 1 h + Gas Normalizing + Solution + Aging 1200 °C, 4 h + Furnace cooling → 1060 °C, 1 h + Gas treatment (NSA) cooling ! 482 C, 4 h + Air cooling treatment (NSA) cooling → 482 °C, 4 h + Air cooling 2.4. Characterizations 2.4. Characterizations Tensile tests were performed on the specimens produced based on the specifications Tensile tests were performed on the specimens produced based on the specifications of round bar specimen no. 3 in the ASTM E8 standard, using a universal testing machine of round bar specimen no. 3 in the ASTM E8 standard, using a universal testing machine (Landmark, MTS, Eden Prairie, MN, USA), with hydraulic wedge grips and an extensome- (Landmark, MTS, Eden Prairie, MN, USA), with hydraulic wedge grips and an extensom- ter. Each specimen had a gauge length of 25 mm and a diameter of 6 mm. Tensile tests eter. Each specimen had a gauge length of 25 mm and a diameter of 6 mm. Tensile tests were performed at room temperature at a specific crosshead speed of 2 mm/min. The were performed at room temperature at a specific crosshead speed of 2 mm/min. The ob- obtained data were analyzed to plot the stress–strain curves and determine the 0.2% proof yield tained d stress, ata were an ultimateal tensile yzed to p strength, lot the and stress elongation –strain curv toes failur and d e. e X-ray termine th diffraction e 0.2% p (XRD) roof yield stress, ultimate tensile strength, and elongation to failure. X-ray diffraction (XRD) analysis was performed using a Cu K source (40 kV and 30 mA) to evaluate the phase analysis was performed using a Cu Kα source (40 kV and 30 mA) to evaluate the phase composition over a 2 range of 20–90 . Microstructural observations were conducted using optical composit micr ion oscopy over a 2 (OM, θ raV nge of ert A1, 20–9 Carl 0°. Zeiss, Microst Oberkoch ructura en, l observat Baden-Württember ions were conduct g, Germany) ed us- and ing op scanning tical mielectr croscopy on micr (OM, oscopy Vert (SEM, A1, Ca JSM-F100, rl Zeiss, Obe JEOL, rkochen, Akishimam, Baden-W Tokyo, ürttemb Japan). erg, Ge The r- specimens for electron backscattered diffraction (EBSD) analysis were polished using ion many) and scanning electron microscopy (SEM, JSM-F100, JEOL, Akishimam, Tokyo, Ja- milling pan). The (cr spe oss-section cimens for polisher electron IB-19520CCP backscattered , JEOL, diffrac Akishimam, tion (EBSD) T aokyo, nalysis Japan) were polishe at 3 kV d for 2 h. EBSD analysis was performed using an SEM and EBSD system (Velocity super, using ion milling (cross-section polisher IB-19520CCP, JEOL, Akishimam, Tokyo, Japan) EDAX, Pleasanton, CA, USA) with an acceleration voltage of 20 kV. To determine copper at 3 kV for 2 h. EBSD analysis was performed using an SEM and EBSD system (Velocity precipitation, a transmission electron microscope (TEM, Talos F200X model, FEI, Hillsboro, super, EDAX, Pleasanton, CA, USA) with an acceleration voltage of 20 kV. To determine OR, USA) was used at 200 kV. Metals 2022, 12, x FOR PEER REVIEW 5 of 14 copper precipitation, a transmission electron microscope (TEM, Talos F200X model, FEI, Metals 2022, 12, 704 5 of 14 Hillsboro, OR, USA) was used at 200 kV. 3. Results and Discussion 3. Results and Discussion 3.1. Mechanical Properties 3.1. Mechanical Properties The effects of the heat treatments on the mechanical properties of L-PBF 17-4 PH SS The effects of the heat treatments on the mechanical properties of L-PBF 17-4 PH SS were evaluated using tensile tests. To obtain reliable and reproducible data, five tests were were evaluated using tensile tests. To obtain reliable and reproducible data, five tests were performed for each HT condition. The results are shown in Figure 4. performed for each HT condition. The results are shown in Figure 4. Figure 4. Yield strength, tensile strength, and elongation for the L-PBF as-built and heat-treated 17-4 Figure 4. Yield strength, tensile strength, and elongation for the L-PBF as-built and heat-treated 17- PH sample. 4 PH sample. The tensile strengths of the AB and SA specimens treated using the conventional heat- The tensile strengths of the AB and SA specimens treated using the conventional ing method were similar to those reported by other researchers, as shown in heating method were similar to those reported by other researchers, as shown in Figure 1 Figure 1 [23,26,27]. Conversely, the tensile strength of the NSA specimen significantly [23,26,27]. Conversely, the tensile strength of the NSA specimen significantly exceeded exceeded that of the 17-4 PH SS manufactured by the L-PBF process by other researchers that of the 17-4 PH SS manufactured by the L-PBF process by other researchers and met and met the requirements of the H900 condition in the ASTM A693 standard. These results the requirements of the H900 condition in the ASTM A693 standard. These results indicate indicate that the mechanical properties were significantly dependent on the application of that the mechanical properties were significantly dependent on the application of normal- normalizing treatment prior to the solution and aging treatments. izing treatment prior to the solution and aging treatments. The yield strength of the AB specimen obtained in this study was 535 MPa, which is The yield strength of the AB specimen obtained in this study was 535 MPa, which is the lowest among those of the specimens obtained under all the HT conditions and lower the lowest among those of the specimens obtained under all the HT conditions and lower than that of the L-PBF 17-4 PH SS produced by other groups. The yield strengths of the S than that of the L-PBF 17-4 PH SS produced by other groups. The yield strengths of the S and SA specimens were 593 MPa and 722 MPa, respectively, which were slightly higher and SA specimens were 593 MPa and 722 MPa, respectively, which were slightly higher than those of the AB specimen but significantly lower than those of the H900 HT-wrought than those of the AB specimen but significantly lower than those of the H900 HT-wrought 17-4 PH SS or the H900 condition of the ASTM A693 standard, as shown in Figure 1. 17-4 PH SS or the H900 condition of the ASTM A693 standard, as shown in Figure 1. Conversely, the yield strength and elongation of the NSA specimen significantly Conversely, the yield strength and elongation of the NSA specimen significantly in- increased to 1265 MPa and 12.95%, respectively, and the standard deviation decreased creased to 1265 MPa and 12.95%, respectively, and the standard deviation decreased com- compared to that of the SA specimen, indicating the uniformity of the mechanical properties pared to that of the SA specimen, indicating the uniformity of the mechanical properties in the NSA specimen. These results indicate that microstructural changes occurred during in the NSA specimen. These results indicate that microstructural changes occurred during the normalizing treatment, affecting the microstructural evolution during the solution and the normalizing treatment, affecting the microstructural evolution during the solution and aging treatments. aging treatments. 3.2. XRD Patterns and Optical Micrography Images XRD analysis was performed to determine the phases and microstructures of the AB and HT specimens. As shown in Figure 5, the AB and N specimens were composed of a large fraction of body-centered cubic (BCC) and certain amount of face-centered cubic (FCC) phases, whereas the S, SA, and NSA specimens consisted almost entirely of the BCC phase. This implies that during cooling after the solution HT, the FCC phase in specimens AB and N was transformed into the BCC phase. Metals 2022, 12, x FOR PEER REVIEW 6 of 14 3.2. XRD Patterns and Optical Micrography Images XRD analysis was performed to determine the phases and microstructures of the AB and HT specimens. As shown in Figure 5, the AB and N specimens were composed of a large fraction of body-centered cubic (BCC) and certain amount of face-centered cubic (FCC) phases, whereas the S, SA, and NSA specimens consisted almost entirely of the BCC phase. This implies that during cooling after the solution HT, the FCC phase in specimens Metals 2022, 12, 704 6 of 14 AB and N was transformed into the BCC phase. Figure 5. XRD patterns of the as-built and heat-treated 17-4 PH samples. Figure 5. XRD patterns of the as-built and heat-treated 17-4 PH samples. The SA and NSA specimens showed no significant differences in their XRD results The SA and NSA specimens showed no significant differences in their XRD results despite exhibiting large differences in their tensile and yield strengths, as shown in Figure 4. despite exhibiting large differences in their tensile and yield strengths, as shown in Figure This could be because the body-centered tetragonal (BCT) martensite and the BCC -ferrite 4. This could be because the body-centered tetragonal (BCT) martensite and the BCC δ- were significantly similar, so there were not distinguishable in the XRD pattern. This is ferrite were significantly similar, so there were not distinguishable in the XRD pattern. because of the fact that the BCT martensite crystal structure in 17-4 PH SS is the BCC phase This is because of the fact that the BCT martensite crystal structure in 17-4 PH SS is the with a small amount of lattice distortion, similar to the distortions seen in significantly BCC phase with a small amount of lattice distortion, similar to the distortions seen in sig- low-carbon SS [14]. nificantly low-carbon SS [14]. To observe the changes in the microstructures, optical micrography images of the To observe the changes in the microstructures, optical micrography images of the cross-sections parallel to the build direction of the AB, N, S, SA, and NSA specimens were cross-sections parallel to the build direction of the AB, N, S, SA, and NSA specimens were examined. The AB specimen exhibited the typical microstructure of L-PBF 17-4 PH SS, with examined. The AB specimen exhibited the typical microstructure of L-PBF 17-4 PH SS, large, elongated grains grown across melt pools along the build direction and equiaxed fine with large, elongated grains grown across melt pools along the build direction and equi- grains at the melt pool boundaries. The large elongated and equiaxed grains in Figure 6a axed fine grains at the melt pool boundaries. The large elongated and equiaxed grains in are reported as -ferrite and retained austenite, respectively [37]. Figure 6a are reported as δ-ferrite and retained austenite, respectively [37]. The phase fraction of the AB L-PBF 17-4 PH SS depended on the Cr /Ni value eq eq of the initial powder [24,25]. The Cr and Ni values were calculated using WRC-1992 eq eq equation [38]: Cr (wt.%) = Cr + Mo + (0.7 Nb) (1) eq Ni (wt.%) = Ni + (35 C) + (20 N) + (0.25 Cu) (2) eq High levels of -ferrite (80–95%) were obtained when the Cr /Ni ratios were be- eq eq tween 2.65 and 2.81, and nearly 25 vol.% of -ferrite appeared when the Cr /Ni ratio eq eq was less than 2.36 [24]. The Cr /Ni value of the powder used in this study was 2.60 eq eq (using WRC-1992 equations), and a relatively high -ferrite fraction was induced. After the heat treatment, the columnar -ferrite, fine martensite, and austenite phases in the AB specimen were reconstructed into equiaxed austenite grains through recovery and recrystallization. When only solution heat treatment was performed, the recrystallized austenite grains were smaller on average and irregular in size, ranging from several micrometers to several tens of micrometers (Figure 6c,d). When the specimens were subjected to normalizing HT maintained at 1200 C for 4 h, austenite grains with sizes Metals 2022, 12, 704 7 of 14 30–50 m were uniformly present owing to grain growth (Figure 6b,e). The Hall–Petch relationship implies that the strength and elongation to failure of a material are enhanced at smaller grain sizes. The optical micrography results suggest that the SA specimen has Metals 2022, 12, x FOR PEER REVIEW 7 of 14 greater strength and elongation than the NSA specimen. However, an opposite result was obtained, as shown in Figure 4. Figure 6. Optical micrographs obtained from etched sections of the as-built and heat-treated 17-4 PH Figure 6. Optical micrographs obtained from etched sections of the as-built and heat-treated 17-4 samples: (a) AB, (b) N, (c) S, (d) SA, (e) NSA. PH samples: (a) AB, (b) N, (c) S, (d) SA, (e) NSA. The normalized specimen included several micrometer-sized spherical particles, which The phase fraction of the AB L-PBF 17-4 PH SS depended on the Creq/Nieq value of were verified to be Nb-rich carbides by EDS (Figure 6b). The Nb-rich carbides in the N the initial powder [24,25]. The Creq and Nieq values were calculated using WRC-1992 equa- specimen became smaller in the NSA specimen matrix because of decomposition after the tion [38]: post-solution and aging treatments (Figure 6e). Creq wt.% CrMo0 .7Nb (1) 3.3. Phase Analysis (2) Nieq wt.% Ni35C20 N0.25Cu Further EBSD analysis was performed to investigate the phases of the AB and HT specimens High levels in detail. of δFigur -ferrite e (80–95% 7 shows a ) were o cross-sectional btained whe view n the Cr of the AB eq/N and ieq ratios were HT specimens be- tween along 2 the .65build and 2. dir 81ection. , and ne The arlyorientation 25 vol.% of and δ-fegrain rrite ap boundaries peared when are the C revealed req/Ni mor eq ra eti clearly o was than those in Figure 6. less than 2.36 [24]. The Creq/Nieq value of the powder used in this study was 2.60 (using The AB specimen microstructure consisted of columnar grains (tens of micrometers WRC-1992 equations), and a relatively high δ-ferrite fraction was induced. in size) and fine-equiaxed grains preferentially located between the columnar grains, as After the heat treatment, the columnar δ-ferrite, fine martensite, and austenite phases shown in Figure 7a. The columnar grains were a fully -ferritic BCC microstructure formed in the AB specimen were reconstructed into equiaxed austenite grains through recovery 5 7 by an extremely high cooling rate of 10 –10 K/s, and the fine-equiaxed grains were a and recrystallization. When only solution heat treatment was performed, the recrystal- mixture of BCC martensitic laths and fine FCC austenite grains. These results are similar to lized austenite grains were smaller on average and irregular in size, ranging from several those obtained in a previous study on an AB 17-4 PH SS fabricated using L-PBF [11]. micrometers to several tens of micrometers (Figure 6c,d). When the specimens were sub- Figure 7b,d shows the microstructures of the specimens after HT. The grain-size jected to normalizing HT maintained at 1200 °C for 4 h, austenite grains with sizes 30–50 differences with and without normalization in the treated specimens are consistent with the μm were uniformly present owing to grain growth (Figure 6b,e). The Hall–Petch relation- results from the optical images. The microstructures of the normalized specimens shown in ship implies that the strength and elongation to failure of a material are enhanced at Figure 7b,e are relatively coarse and homogeneous, whereas those without normalization smaller grain sizes. The optical micrography results suggest that the SA specimen has (Figure 7c,d) are smaller and contain some coarse grains. greater strength and elongation than the NSA specimen. However, an opposite result was obtained, as shown in Figure 4. The normalized specimen included several micrometer-sized spherical particles, which were verified to be Nb-rich carbides by EDS (Figure 6b). The Nb-rich carbides in the N specimen became smaller in the NSA specimen matrix because of decomposition after the post-solution and aging treatments (Figure 6e). 3.3. Phase Analysis Further EBSD analysis was performed to investigate the phases of the AB and HT specimens in detail. Figure 7 shows a cross-sectional view of the AB and HT specimens along the build direction. The orientation and grain boundaries are revealed more clearly than those in Figure 6. Metals 2022, 12, x FOR PEER REVIEW 8 of 14 Metals 2022, 12, 704 8 of 14 Figure 7. EBSD orientation maps obtained from the as-built and heat-treated 17-4 PH samples: (a) Figure 7. EBSD orientation maps obtained from the as-built and heat-treated 17-4 PH samples: AB, (b) N, (c) S, (d) SA, (e) NSA. (a) AB, (b) N, (c) S, (d) SA, (e) NSA. The AB specimen microstructure consisted of columnar grains (tens of micrometers The phase fractions of all the post-HT specimens were similar, consisting mostly of the in size) and fine-equiaxed grains preferentially located between the columnar grains, as ferrite phase and approximately 1% austenite, except for the N specimen, which contained shown in approximately Figure 10% 7a.austenite The columnar (each phase grains map wer and e a fraction fully δ-fof errit allic BCC microst specimens is shown ructure in 5 7 formed Supplementary by an extreme Figur le y hig S1 and h cooling Table r S2). ate of This 10 r –10 esult K/ is s, a contradictory nd the fine-eq to uithat axed gra of a pr ins were evious study, in which the incomplete transformation of martensite occurred more easily when a mixture of BCC martensitic laths and fine FCC austenite grains. These results are similar to nitr those ogen obt gas aiwas ned in used a previ instead ous of study argon on gas an A asB the 17build-chamber -4 PH SS fabrica envir ted u onment sing L-P [12 BF [ ]. In 11]this . study, gas cooling was used instead of air cooling in the solution treatment process to avoid Figure 7b,d shows the microstructures of the specimens after HT. The grain-size dif- the retained austenite phase and obtain a full martensite phase. As expected, a significantly ferences with and without normalization in the treated specimens are consistent with the small fraction of the austenite phase was found in the specimens treated by gas cooling, as results from the optical images. The microstructures of the normalized specimens shown shown in Figure 7c,e. in Figure 7b,e are relatively coarse and homogeneous, whereas those without normaliza- The EBSD results shown in Figure 7d,e cannot explain the difference between the tion (Figure 7c,d) are smaller and contain some coarse grains. mechanical properties of the SA and NSA specimens, as shown in Figure 4. Additional The phase fractions of all the post-HT specimens were similar, consisting mostly of analyses were performed to identify the cause and mechanism for this result. the ferrite phase and approximately 1% austenite, except for the N specimen, which con- tained approximately 10% austenite. This result is contradictory to that of a previous 3.4. Cu Precipitation study, in which the incomplete transformation of martensite occurred more easily when Age hardening is an HT process for increasing the yield strength through the uniform nitrogen gas was used instead of argon gas as the build-chamber environment [12]. In this precipitation of 10–20 nm Cu particles and inhibiting the dislocation movement through study, gas cooling was used instead of air cooling in the solution treatment process to the Orowan mechanism. The yield strengths of the 17-4 PH SS specimens were significantly avoid the retained austenite phase and obtain a full martensite phase. As expected, a sig- affected by the amount of Cu precipitation. Therefore, TEM and EDS mapping analyses nificantly small fraction of the austenite phase was found in the specimens treated by gas were performed to verify the correlation between the microstructure and the mechanical cooling, as shown in Figure 7c,e. properties of the SA and NSA specimens after aging, as shown in Figure 8. The EBSD results shown in Figure 7d,e cannot explain the difference between the Nanoscale Cu-rich precipitates were observed in both the SA (Figure 8a) and the mechanical properties of the SA and NSA specimens, as shown in Figure 4. Additional NSA (Figure 8d) specimens (results of the TEM-EDS map about the distribution of major analyses were performed to identify the cause and mechanism for this result. compositions are shown in Figures S2–S6). The SA specimen displayed two distinct types of regions: Figure 8b shows no precipitation or only precipitates inhomogeneously, and 3.4. Cu Precipitation Figure 8c shows Cu particles precipitated homogeneously. Figure 8a was magnified by Age hardening is an HT process for increasing the yield strength through the uniform 2.5 times in Figure 8b,c to clearly show the differences between the two regions. Ho- precipitation of 10–20 nm Cu particles and inhibiting the dislocation movement through mogeneous Cu precipitation was obtained after the normalizing treatment, as shown the Orowan mechanism. The yield strengths of the 17-4 PH SS specimens were signifi- in Figure 8d,e. cantly affected by the amount of Cu precipitation. Therefore, TEM and EDS mapping anal- yses were performed to verify the correlation between the microstructure and the me- chanical properties of the SA and NSA specimens after aging, as shown in Figure 8. Metals 2022, 12, x FOR PEER REVIEW 9 of 14 Metals 2022, 12, 704 9 of 14 Figure Figure 8. 8. Com Comparison parison of T of TEM-E EM-EDS analysis o DS analysis of f 17-4 PH SS 17-4 PH SS sh showing owing nano-sc nano-scale ale Cu Cu-rich -rich pre precipitation. cipitation. These results explain the difference in the yield strengths between the SA and the NSA Nanoscale Cu-rich precipitates were observed in both the SA (Figure 8a) and the NSA specimens, as shown in Figure 4. The conditions required to form the ideal precipitate Cu (Figure 8d) specimens. The SA specimen displayed two distinct types of regions: Figure particle distribution are considered below. 8b shows no precipitation or only precipitates inhomogeneously, and Figure 8c shows Cu particles precipitated homogeneously. Figure 8a was magnified by 2.5 times in Figure 8b, 3.5. Effects of Normalizing Treatment c to clearly show the differences between the two regions. Homogeneous Cu precipitation Figure 9 shows the equilibrium phase diagram of 17-4 PH SS simulated using JMatPro was obtained after the normalizing treatment, as shown in Figure 8d,e. software. The results show that the solidification of liquid 17-4 SS begins through the These results explain the difference in the yield strengths between the SA and the nucleation of primary -ferrites at approximately 1430 C and that the -ferrite phase NSA specimens, as shown in Figure 4. The conditions required to form the ideal precipi- transforms into the austenite phase by solid-state diffusion as the temperature decreases tate Cu particle distribution are considered below. from 1320 to 1120 C. It is well known that the phase transformation of austenite into martensite by rapid cooling occurs at the starting temperature Ms (132 C) and ends at the 3.5. Effects of Normalizing Treatment temperature Mf (32 C) of martensitic transformation [39]. The final structure formed along Figure 9 shows the equilibrium phase diagram of 17-4 PH SS simulated using these paths would then be mostly martensite. JMatPro software. The results show that the solidification of liquid 17-4 SS begins through The L-PBF fabrication process of 17-4 PH SS is different and follows the paths of the nucleation of primary δ-ferrites at approximately 1430 °C and that the δ-ferrite phase the non-equilibrium solidification processes; thus, it does not agree with the equilibrium transforms into the austenite phase by solid-state diffusion as the temperature decreases state. The simulation results obtained in the equilibrium state predict the solidification from 1320 to 1120 °C. It is well known that the phase transformation of austenite into and cooling processes and cannot explain the kinetics of the governing phase transforma- martensite by rapid cooling occurs at the starting temperature Ms (132 °C) and ends at the tions [24]. At the end of the non-equilibrium solidification process, the non-equilibrium temperature Mf (32 °C) of martensitic transformation [39]. The final structure formed microstructures are still highly likely to exist in the final post-HT state, unless sufficient time along these paths would then be mostly martensite. and energy are provided at the post-HT stage. Therefore, a kernel average misorientation (KAM) analysis was performed to identify the remaining non-equilibrium phases. Figure 10 shows the KAM graphs of the SA (Figure 10a) and NSA (Figure 10b) speci- mens. The KAM graphs quantify the average local misorientation of the nearest neighbors around the measurement point, from which the lattice distortion and dislocation accu- mulation can be deduced. The results in the KAM graphs are consistent with the larger lattice distortion in the martensite and the smaller distortion in the ferrite microstructure. It was previously reported that the martensite structure of 17-4 PH SS has a high disloca- Metals 2022, 12, 704 10 of 14 Metals 2022, 12, x FOR PEER REVIEW 10 of 14 tion density, whereas ferrite has a comparatively small dislocation density after solution treatment [17,19]. Metals 2022, 12, x FOR PEER REVIEW 11 of 14 Figure 9. Equilibrium phase diagram of 17-4 PH SS calculated by JMatPro software. Figure 9. Equilibrium phase diagram of 17-4 PH SS calculated by JMatPro software. The L-PBF fabrication process of 17-4 PH SS is different and follows the paths of the non-equilibrium solidification processes; thus, it does not agree with the equilibrium state. The simulation results obtained in the equilibrium state predict the solidification and cool- ing processes and cannot explain the kinetics of the governing phase transformations [24]. At the end of the non-equilibrium solidification process, the non-equilibrium microstruc- tures are still highly likely to exist in the final post-HT state, unless sufficient time and energy are provided at the post-HT stage. Therefore, a kernel average misorientation (KAM) analysis was performed to identify the remaining non-equilibrium phases. Figure 10 shows the KAM graphs of the SA (Figure 10a) and NSA (Figure 10b) spec- imens. The KAM graphs quantify the average local misorientation of the nearest neigh- bors around the measurement point, from which the lattice distortion and dislocation ac- cumulation can be deduced. The results in the KAM graphs are consistent with the larger lattice distortion in the martensite and the smaller distortion in the ferrite microstructure. It was previously reported that the martensite structure of 17-4 PH SS has a high disloca- tion density, whereas ferrite has a comparatively small dislocation density after solution treatment [17,19]. Comparing the SA and NSA specimens, Figure 10a is largely blue (indicating little Figure 10. Kernel average misorientation graphs of the 17-4 PH steel heat treated by (a) SA and Figure 10. Kernel average misorientation graphs of the 17-4 PH steel heat treated by (a) SA and (b) lattice distortion), and Figure 10b shows larger green and red areas. This means that the (b) NSA. NSA. SA specimen HT with only the solution process and without the normalizing treatment Comparing the SA and NSA specimens, Figure 10a is largely blue (indicating little had a higher probability of including the ferrite phase and not the martensite phase. Con- The authors suggested that the presence of the ferrite phase after solution HT was lattice distortion), and Figure 10b shows larger green and red areas. This means that the SA versely, the NSA specimen subjected to both normalizing and solution treatments con- caused by insufficient transformation into the solid state owing to the low diffusion rate specimen HT with only the solution process and without the normalizing treatment had a tained more martensite phases. at 1040 °C [37,40]. Moreover, the 17-4 PH SS powder used in this study had a high Cr higher probability of including the ferrite phase and not the martensite phase. Conversely, the NSA specimen subjected to both normalizing and solution treatments contained more content of 17.10%. Cr acted as a ferrite-stabilizing element, suggesting that more time may martensite phases. be required for solid-state diffusion to transform ferrite into austenite. Nezhadfar et al. The authors suggested that the presence of the ferrite phase after solution HT was [40] confirmed the presence of δ-ferrite after solution HT (1050 °C) and aging HT (482 °C) caused by insufficient transformation into the solid state owing to the low diffusion rate of L-PBF 17-4 PH SS with high Creq/Nieq values. at 1040 C [37,40]. Moreover, the 17-4 PH SS powder used in this study had a high Cr Various factors affect Cu precipitation behavior, such as the matrix (δ-ferrite, mar- tensite, and austenite) and microsegregation (ferrite, austenite stabilizer, and precipitation elements). Many researchers have reported that ferrite remains even after HT and that δ- ferrite in 17-4 PH SS has various detrimental effects on hardening behavior [4,17,41]. Sabooni et al. [17] reported that Cu precipitation is delayed in 17-4 PH with a ferrite structure compared to 17-4 PH with a martensite structure. During prolonged aging of the martensite structure, decreases in strength and hardness may occur prematurely through the recovery of dislocations, the coarsening of the Cu-rich precipitates, and the formation of reverted austenite. Additionally, the high dislocation density in the marten- sitic structure accelerates the precipitation of Cu particles by facilitating pipe diffusion and providing heterogeneous nucleation sites [19]. As shown in Section 3.3, the grain size of 17-4 PH SS fabricated by L-PBF was in- creased by the normalizing treatment through grain growth and the promotion of Cu pre- cipitation as the ferrite dissolved. It has been reported that the extremely fine-grain size that frequently appears in the AM process lowers the Ms temperature and inhibits the martensitic transformation [42–44]. Colaço et al. [21] reported that the grain size was re- duced from 100 to 2.7 μm through the laser surface melting of AISI 420 martensitic SS, and the martensitic transformation temperature decreased from 115 to 20 °C. Therefore, it can be concluded that the normalizing treatment of L-PBF-fabricated 17- 4PH SS facilitates full martensite transformation by dissolving ferrite and triggering grain growth and thereby enabling homogeneous Cu precipitation during the aging HT. There- fore, the mechanical properties of 17-4 PH SS fabricated using L-PBF in a nitrogen gas environment can be improved by adding a normalizing treatment to the conventional process. 4. Conclusions In this study, the effects of the HT profile on the microstructure and mechanical prop- erties of 17-4 PH SS fabricated by L-PBF were studied, and the following conclusions were drawn from the results of a comprehensive series of investigations: Metals 2022, 12, 704 11 of 14 content of 17.10%. Cr acted as a ferrite-stabilizing element, suggesting that more time may be required for solid-state diffusion to transform ferrite into austenite. Nezhadfar et al. [40] confirmed the presence of -ferrite after solution HT (1050 C) and aging HT (482 C) of L-PBF 17-4 PH SS with high Creq/Nieq values. Various factors affect Cu precipitation behavior, such as the matrix (-ferrite, marten- site, and austenite) and microsegregation (ferrite, austenite stabilizer, and precipitation elements). Many researchers have reported that ferrite remains even after HT and that -ferrite in 17-4 PH SS has various detrimental effects on hardening behavior [4,17,41]. Sabooni et al. [17] reported that Cu precipitation is delayed in 17-4 PH with a ferrite structure compared to 17-4 PH with a martensite structure. During prolonged aging of the martensite structure, decreases in strength and hardness may occur prematurely through the recovery of dislocations, the coarsening of the Cu-rich precipitates, and the formation of reverted austenite. Additionally, the high dislocation density in the martensitic structure accelerates the precipitation of Cu particles by facilitating pipe diffusion and providing heterogeneous nucleation sites [19]. As shown in Section 3.3, the grain size of 17-4 PH SS fabricated by L-PBF was increased by the normalizing treatment through grain growth and the promotion of Cu precipitation as the ferrite dissolved. It has been reported that the extremely fine-grain size that fre- quently appears in the AM process lowers the Ms temperature and inhibits the martensitic transformation [42–44]. Colaço et al. [21] reported that the grain size was reduced from 100 to 2.7 m through the laser surface melting of AISI 420 martensitic SS, and the marten- sitic transformation temperature decreased from 115 to 20 C. Therefore, it can be concluded that the normalizing treatment of L-PBF-fabricated 17-4PH SS facilitates full martensite transformation by dissolving ferrite and trigger- ing grain growth and thereby enabling homogeneous Cu precipitation during the ag- ing HT. Therefore, the mechanical properties of 17-4 PH SS fabricated using L-PBF in a nitrogen gas environment can be improved by adding a normalizing treatment to the conventional process. 4. Conclusions In this study, the effects of the HT profile on the microstructure and mechanical properties of 17-4 PH SS fabricated by L-PBF were studied, and the following conclusions were drawn from the results of a comprehensive series of investigations: Significant differences in the mechanical properties of 17-4 PH SS were observed among the various HTs. Compared to the conventional method, the NSA treatment condition resulted in superior mechanical properties that met the ASTM-693 (H900) industrial standard. The yield strength and elongation of the resulting 17-4 PH SS were 1264 MPa and 12.9%, respectively. No differences in the austenite fraction were found in the XRD and EBSD results after the additional normalizing treatment. However, irregular fine-grain size and inhomo- geneous distribution of the Cu precipitates that affected the mechanical properties of the 17-4 PH SS under SA conditions were observed in the OM and TEM-EDS analyses. The existence of a non-equilibrium ferrite phase under the SA condition was observed using KAM, and it was confirmed that the mechanical properties changed based on the differences in the precipitation behavior of the ferrite and martensite structures. Cu precipitation in 17-4 PH SS fabricated by L-PBF in a nitrogen environment can be stabilized by applying an additional normalizing treatment. Inhomogeneous Cu precipitation occurs in 17-4 PH SS without a normalizing step because of the mixed ferrite and martensite microstructure. Metals 2022, 12, 704 12 of 14 Supplementary Materials: The following are available online at https://www.mdpi.com/article/10 .3390/met12050704/s1, Figure S1: EBSD orientation maps and Phase maps (green: ferrite and/or martensite, red: austenite) obtained from the as-built and heat treated 17-4 PH samples: (a-b) AB, (c-d) N, (e-f) S, (g-h) SA, (i-j) NSA; Figure S2: TEM-EDS analysis of SA specimen showing inhomoge- neous Cu-rich precipitation (magnification: 450 kx); Figure S3: TEM-EDS analysis of SA specimen showing nano-scale Cu-rich precipitation (magnification: 910 kx); Figure S4: TEM-EDS analysis of SA specimen showing homogeneous Cu-rich precipitation (magnification: 910 kx); Figure S5: TEM-EDS analysis of NSA specimen showing homogeneous nano-scale Cu-rich precipitation (magnification: 450 kx); Figure S6: TEM-EDS analysis of NSA specimen showing homogeneous nano-scale Cu-rich precipitation (magnification: 910 kx); Table S1: Mechanical properties of 17-4 PH SS produced by L- PBF from previous studies [23,26–31]; Table S2: Result of EBSD-Phases map showing phases fraction of AB and HT specimens. Author Contributions: S.-M.Y. and J.Y. wrote and edited the manuscript and contributed to all activities. T.B.K. and S.H.L. fabricated STS 630 parts and conducted experiments and heat treatment. K.C., T.-S.J. and Y.S. contributed to the interpretation and discussion of the experimental results. All authors have read and agreed to the published version of the manuscript. Funding: This research received no external funding. Institutional Review Board Statement: Not applicable. Informed Consent Statement: Not applicable. Data Availability Statement: The data presented in this study are available upon request from the corresponding author. Acknowledgments: This research was performed through R&D projects (Grant No. 20009815 and 20013122), financially supported by the Ministry of Trade, Industry, and Energy (MOTIE) and korea institute of industrial technology internal project (Grant No. EH220013) in Korea. 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Metals – Multidisciplinary Digital Publishing Institute
Published: Apr 20, 2022
Keywords: 17-4 PH stainless steel; selective laser melting; laser powder bed fusion; heat treatment; microstructure; mechanical properties
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