Effect of scanning speed on the microstructure and mechanical behavior of 316L stainless steel fabricated by selective laser melting

Journal Pre-proof Effect of scanning speed on the microstructure and mechanical behavior of 316L stainless steel fabricated by selective laser melting

Jiangwei Liu, Yanan Song, Chaoyue Chen, Xiebin Wang, Hu Li, Chang'an Zhou, Jiang Wang, Kai Guo, Jie Sun PII:

S0264-1275(19)30793-2

DOI:

https://doi.org/10.1016/j.matdes.2019.108355

Reference:

JMADE 108355

To appear in:

Materials & Design

Received date:

14 August 2019

Revised date:

19 October 2019

Accepted date:

12 November 2019

Please cite this article as: J. Liu, Y. Song, C. Chen, et al., Effect of scanning speed on the microstructure and mechanical behavior of 316L stainless steel fabricated by selective laser melting, Materials & Design(2018), https://doi.org/10.1016/j.matdes.2019.108355

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© 2018 Published by Elsevier.

Journal Pre-proof Effect of scanning speed on the microstructure and mechanical behavior of 316L stainless steel fabricated by selective laser melting Jiangwei Liua,b, Yanan Songa,b, Chaoyue Chenc, Xiebin Wangd, Hu Lie, Chang’an Zhoua,b, Jiang Wangc, Kai Guoa,b,*, Jie Suna,b a

Key Laboratory of High Efficiency and Clean Mechanical Manufacture, School of Mechanical Engineering, Shandong University, Jinan 250061, P.R. China

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Research Centre for Aeronautical Component Manufacturing Technology and Equipment, Shandong

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University, Jinan 250061, P.R. China State Key Laboratory of Advanced Special Steels, School of Materials Science and Engineering,

School of Materials Science and Engineering, Shandong University, Jinan 250061, P.R. China

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Shanghai University, Shanghai 200444, P.R. China

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School of Electrical and Electronic Engineering, University of Manchester, M13 9PL Manchester, United Kingdom

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Abstract: In this work, the tensile behavior and microhardness of 316L stainless steel fabricated by selective laser melting under different process parameters were investigated. The ultimate tensile

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strength decreased slightly with increasing energy input, while the opposite tendency was observed for the elongation to failure. Microstructure characterizations were performed to relate the pore

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morphology, melting pool geometry, solidification cell structure, and grain sizes to the mechanical performance of as-built samples. Fine grains with high fraction of low-angle grain boundaries and fine cellular structures with nano-inclusions are observed in the sample fabricated with a high scanning speed (1000 mm/s). As a result, the sample shows high ultimate tensile strength of up to 707 MPa, while maintaining a total elongation of 30%. The sample fabricated with a low scanning speed (800 mm/s) shows high ductility with total elongation to failure of 55%. The improved ductility is mainly attributed to the elimination of residual pores and melting pool boundaries, which result in brittle features in the as-built samples. The results indicate that selective laser melting may act as a physical metallurgy method to modify the microstructure, and thus improve the mechanical performance of metallic materials. Keywords: Selective laser melting; 316L stainless steel; Scanning speed; Mechanical performance;

Journal Pre-proof Microstructure characterization 1. Introduction Austenitic stainless steel is widely used as an engineering material in various industries owing to its excellent corrosion and oxidation resistance, as well as its good formability [1–4]. The presence of coarse grains that cause low tensile strength is the main factor that hinders the industrial applications of 316L stainless steel. Therefore, manufacturing strong and ductile stainless steel has become a central research topic in recent years. Conventional ways to strengthen stainless steel require

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additional tooling or treatment after fabrication [5–9]; however, this is not practical owing to the limitations in manufacturing components with complex geometry.

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Additive manufacturing (AM) plays an increasing role in modern industry and is considered as an

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alternative method that offers the possibility of producing parts with densified structures and

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improved mechanical properties from various metals [10–14]. Selective laser melting (SLM) is one of the most widely used AM techniques, which can locally melt pre-deposited powders using a

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high-energy laser beam. The fabrication of 316L stainless steel by SLM to obtain specimens with high strength and ductility has been reported in the literature [15,16]. It was found that the

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manufacturing orientation greatly affected the mechanical properties: the yield strength and tensile strength of specimens built in the horizontal direction were usually higher than those built in the

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perpendicular direction (built-up direction) [17]. Zhang et al. found that the comprehensive optimization of placement orientation and scanning angle led to highly densified 316L stainless steel with the highest tensile strength of 657 MPa [18]. Extensive work to relate the process parameters to the microstructure and mechanical properties of stainless steel was conducted by Zhang et al. [19]. They found that yield and tensile strength decreased with increasing laser power or decreasing scanning speed; however, there was no consistent tendency in elongation with changes in laser power or scanning speed. Furthermore, it has been proved that the melting pool boundary is of great importance for SLM-manufactured 316L stainless steel [20]. Further investigations need to be carried out to relate the process parameters to the melting pool geometry, microstructure, and mechanical performance of SLM-manufactured 316L stainless steel. In this work, a set of processing parameters was used to fabricate 316L stainless steel based on SLM.

Journal Pre-proof Tensile tests and microhardness measurements were performed to determine the effects of process parameters on the mechanical performance of the as-built samples. The porosity, melting pool morphology, grain size, and crystallographic orientation were subsequently evaluated to relate the microstructure to the mechanical properties of the as-built 316L stainless steel. 2. Material and methods 2.1. SLM process The feedstock material used in this work was austenitic stainless-steel powder with a grade of

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LaserForm 316L (A) (3D Systems Inc., USA). Fig. 1a illustrates the scanning electron microscopy (SEM) overview of the spherical powder. Fig. 1b shows the particle size distribution of the feedstock

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powder, ranging from 6 μm to 58 μm.

Fig. 1 (a) SEM overview and (b) particle size distribution of the feedstock powder. A commercial selective laser melting machine (ProX DMP 320, 3D Systems Inc., USA) was employed to fabricate 316L stainless steel. A fiber laser with a focused diameter of 80 μm was installed in a sealed chamber filled with ultra-pure argon (>99.99%). Plate-shaped samples (60 mm × 10 mm × 3 mm) were fabricated with the building direction (BD) along the sample width, as shown in Fig. 2a. The scanning strategies used during fabrication are illustrated in Fig. 2b. In each layer, stripe-type scanning was performed with the stripe length of 13 mm, and the laser beam followed bi-directional scanning. The rotation angle of 115° was implemented between successive layers. This angle is one optimum-spaced number between 1° and 360° that can fill every 5° when returning to the initial angle. The layer thickness and hatch spacing were 30 μm and 100 μm, respectively. For the SLM process, the parameters influencing the energy input are laser power (P), scanning speed (V),

Journal Pre-proof layer thickness (τ), and hatch spacing (h) [21]. The layer thickness and hatch spacing were constant, while the energy input per unit volume (E) was defined to evaluate the energy input by combining the contributions of the laser power and scanning speed: 𝐸=

𝑃 𝑉∙𝜏∙ℎ

(1)

The process parameters are summarized in Table 1 to evaluate the effect of laser power (P1, P2, and

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P3) and scanning speed (V1, V2, and V3) on the mechanical performance of the as-built samples.

Fig. 2 (a) Schematic of selective laser melting process, and (b) scanning strategy used for fabricating

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316 L stainless steel samples. Table 1 Detailed process parameters used to fabricate 316L stainless steel samples.

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P1 P2 P3 V1 V2 V3

Laser power (W) 185 215 245 215 215 215

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Case

Scanning speed (mm/s) 900 900 900 800 900 1000

Energy density (J/mm3) 83.6 97.1 110.7 109.2 97.1 87.4

2.2. Mechanical tests The tensile properties of the as-built 316L stainless steel samples were tested by employing a universal testing machine at room temperature with a constant strain rate of 0.02/min. Dog-bone-shaped tensile specimens were machined from the fabricated plates (Fig. 2a). The gauge length, width, and thickness of the tensile specimens were 20 mm, 4 mm, and 2 mm, respectively. The tensile axis of each specimen was along the sample length, which was perpendicular to the building direction. On the stress-strain curve, the ultimate tensile strength (UTS) was taken to evaluate the strength of as-built 316L stainless steel samples, and the tensile elongation to failure was

Journal Pre-proof used to estimate the ductility. The microhardness was measured by a Vickers hardness tester at a load of 0.2 kg and dwell time of 15 s. Each demonstrated microhardness value was the average of five tests. 2.3. Specimen characterization The formed phase of the as-built 316L stainless steel was determined by an X-ray diffraction (XRD) spectrometer (Rigaku Ultima IV) using Cu-Kα radiation from 10° to 90° with a step size of 0.02°. An X-ray computer tomography (X-CT) system (X5000, North Star Imaging) was employed to

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characterize the pore morphology and distribution. In this work, the SLM-built specimens used for X-CT were machined into cylinders with diameter of 2.8 mm. The 3D reconstruction of individual

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pores was realized under the resolution of 4 μm/voxel [22]. SEM (FEI Quanta FEG 250) was

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employed to characterize the microstructure and fracture surface. The grain size and crystallographic

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orientation of the as-built samples were investigated using electron backscatter diffraction (EBSD), which was performed on a field emission gun SEM machine (JEOL JSM-7800F) equipped with an

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Oxford NordlysMax3 system. Considering the scanning efficiency and resolution, the step size of 1 μm was used for EBSD measurements [23]. Prior to the structure characterization, all samples were

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sectioned parallel to the building direction. After the necessary grinding and polishing, the samples

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were etched using a solution of 0.1 mol/L oxalic acid at 6 V. 3. Results and discussions

3.1. Mechanical performance

Fig. 3 reveals that the SLM-manufactured 316L stainless steel has high UTS, much higher than that of as-casted (450 MPa [24]) and wrought samples (560 MPa [25]). For the ductility, total elongation to failure of as-built samples in this work is basically higher than the wrought samples (30% [25]), and comparable to the as-casted ones (45% [24]). for the As shown in Fig. 3a, the UTS decreased slightly with increasing energy density, and the maximum UTS reached 707 MPa when applying the laser power of 215 W and the scanning speed of 1000 mm/s (case V3 in Table 1). In contrast, the elongation to failure increased slightly while increasing the input energy (Fig. 3b). The maximum elongation to failure of 55% occurred with the laser power of 215 W and the scanning speed of 800 mm/s. The high deviation observed for elongation to failure could be attributed to unexpected defects

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during SLM fabrication.

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Fig. 3 Tensile properties of as-built samples with variations of energy density. The values of UTS and elongation to failure are summarized in Fig. 4 for comparison with values in

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literature [15,17,18,24,26–30]. The reported UTS in this work is within the range of 625–707 MPa,

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while the elongation to failure varies from 30% to 55%. High strength-ductility combination was

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achieved in this work.

Fig. 4 Summary of the ultimate tensile strength versus the elongation to failure of SLM-built 316L stainless steel. The Vickers microhardness of the as-built samples is shown in Fig. 5. It shows a progressively increasing trend with variations in energy density. As reported in the literature, for SLM-built 316L stainless steel, the hardness tends to increase linearly with energy density between 48 and 125 J/mm3 [31], and then declines steadily with further increase in energy density [26,32]. In this work, the corresponding energy input per unit volume was on the range of 84–111 J/mm3, suggesting that the increase in microhardness with energy density is reasonable. The microhardness ranged from 207

Journal Pre-proof HV0.2 to 219 HV0.2, which is considerably higher than that of the annealed stainless steel (~150 HV). It is worth noting that the microhardness obtained was comparable to the reported values in the literature [21,26,27,31]. Although it is believed the porosity, microstructure, and oxygen content could have a significant influence on microhardness [31], the factor contributing the most to microhardness remains

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unknown [32].

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Fig. 5 Vickers microhardness under different energy densities.

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For stainless steel, microhardness testing is an efficient means of assessing mechanical properties, and the yield strength increases linearly by increasing the microhardness [33]. With regard to the

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mechanical properties of additive manufactured 316L stainless steel, Smith et al. found that the increase in energy density resulted in the increasing of the yield strength [16], while the opposite

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tendency for the tensile strength was reported by Ma et al. [34]. Therefore, the tendency of mechanical properties of SLM-built 316L stainless steel is highly dependent on the processing parameters, as well as the porosity and microstructure of the as-built samples [31]. Depending on the mechanical performance of the as-built samples, it was found that the maximum UTS occurred with laser power of 215 W and scanning speed of 1000 mm/s (case V3). The maximum elongation to failure was achieved with laser power of 215 W and scanning speed of 800 mm/s (case V1). Therefore, it is of great interest to explore the influence of the process parameters used in the two cases on the measured mechanical properties by investigating the fracture mechanism, microstructure, and crystallographic structure of the SLM-built samples. The morphology of the tensile fracture surface was characterized to understand the failure mechanism during the tensile test. Fig. 6a and 6b present the fracture surface with scanning speed of 800 mm/s, in which the maximum elongation to failure was achieved. Fig. 6c and 6d show the

Journal Pre-proof fracture surface with scanning speed of 1000 mm/s, in which the maximum UTS is obtained. The undesirable crater-like voids shown in Fig. 6 usually act as the main sites that trigger fracture by crack nucleation and propagation during tensile testing, resulting in localized brittleness surrounding those voids [26]. In addition, the melting pool boundaries between successive tracks may induce crack initiation when tensile loading increases to a certain extent [27], leading to brittle features shown in Fig. 6. Similar brittle features were reported as the “cleavage steps” [20,35] or defined as the “quasi-cleavage planes” [7]. They are considered as an important contributor to the tensile fracture of

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SLM samples. Tearing ridges in Fig. 6 were reported as the consequence of the interaction between the crack propagation and grain surfaced plastic deformation [5]. Ductile dimple fracture was also

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observed in both cases, indicating that the fracture mechanism is a combination of ductile and brittle

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mechanism at different scanning speeds.

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rupture. However, more information should be given to evaluate the dominant factor of the fracture

Journal Pre-proof Fig. 6 SEM morphology of the tensile fracture surfaces at different locations: scanning speeds of 800 mm/s (a, b) and 1000 mm/s (c, d). 3.2. Structure characterization Fig. 7 shows the XRD patterns of samples fabricated at different scanning speeds (800 mm/s, 900 mm/s, and 1000 mm/s). The diffraction peaks and intensity magnitudes were similar for the three cases, while the diffraction peaks originated from the γ-phase. Similar characteristic peaks were also confirmed in other literature [3,19,26]. It can be deduced from the XRD patterns that the

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solidification and high cooling rate during SLM [19,36].

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microstructure consists of mono-phase γ resulting from the components of as-built samples, rapid

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Fig. 7 XRD patterns of SLM-manufactured 316L stainless steel with different scanning speeds. Fig. 8 shows 3D reconstructions of X-CT images of the three cylinders at different scanning speeds. The colored ellipsoidal- and spherical-shaped pores were detected inside the as-built specimens (outlined as a grey cylinder). It is evident that optimal scanning speed resulted in the microstructure with barely any pores (Fig. 8b), while the defect density increased at high scanning speed of 1000 mm/s (Fig. 8c). The statistical analysis of the X-CT results provided pore characterizations in terms of porosity, size, and shape, which are summarized in Table 2. The sphericity was used to evaluate how spherical a pore is, and the detailed definition was provided in the literature [37]. The aspect ratio was the ratio of the length to width of one pore. From the porosity shown in Table 2, the density of as-built 316L stainless steel achieved 99.9% despite the variation of scanning speeds, while the density was 99.99% at lower scanning speed. The mean sphericity of pores in cases V1 and V2 was 0.99, while that of case V3 was 0.94. This indicates that the pores within the as-built samples were

Journal Pre-proof mostly spherical rather than flat in shape. The mean aspect ratio of pores was 1.73 for case V1, with the minimum and maximum values ranging from 1.16 to 3.69, respectively. For case V3, the mean aspect ratio was 1.87, demonstrating the high frequency of ellipsoidal pores inside the as-built specimen when applying high scanning speed. Moreover, the equivalent diameter reached 16 μm (ranging from 10 μm to 80 μm) at increased scanning speed of 1000 mm/s. It appears that relatively

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large and ellipsoidal pores may form owing to incomplete melting at high scanning speed.

Fig. 8 Residual pore distributions measured with X-CT: scanning speeds of (a) 800 mm/s, (b) 900

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mm/s, and (c) 1000 mm/s. Table 2 Pore characterizations in terms of porosity, size, and shape as a function of scanning speeds.

V1

0.002

V2

0.002

V3

0.013

Mean sphericity (min./max.) 0.99 ± 0.04 (0.74/1.11) 0.99 ± 0.04 (0.77/1.10) 0.94 ± 0.07 (0.55/1.07)

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Porosity (%)

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Case

Mean aspect ratio (min./max.) 1.73 ± 0.37 (1.16/3.69) 1.79 ± 0.35 (1.16/2.91) 1.87 ± 0.45 (1.18/4.83)

Equivalent diameter (mm) (min./max.) 0.013 ± 0.004 (0.008/0.035) 0.012 ± 0.004 (0.008/0.038) 0.016 ± 0.006 (0.010/0.080)

To further characterize the pores shown in Fig. 8c (case V3), Fig. 9 illustrates the detailed distribution of the equivalent diameter, sphericity, and aspect ratio inside the as-built sample. The majority (>80%) of pore diameters ranged from 10 μm to 20 μm (Fig. 9a). Within this range, the sphericity was above 0.9 (Fig. 9b), and the aspect ratio was below 2 (Fig. 9c). As for the pores with an equivalent diameter of 70–80 μm, the sphericity and aspect ratio were 0.6 and 3, respectively. In addition, the sphericity decreased continuously with the increase in equivalent diameter, while the aspect ratio increased sharply when the equivalent diameter was over 45 μm. In general, the small/ellipsoidal pores were the representative defects, while the large/flat defects were not common

Journal Pre-proof according to the statistical distribution (Fig. 9a), but quite visible owing to their size (as shown in Fig.

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8c).

Fig. 9 (a) Distribution of equivalent diameter, (b) sphericity, and (c) aspect ratio at scanning speed of

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1000 mm/s.

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Spherical pores were generally reported as a result of gas entrapment or the keyhole effect [21].

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Considering the insufficient dwell time of melting pool as well as insufficient energy input when applying high scanning speed, the gas between powder particles could not be entirely released, and

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thus was stuck inside the melting pool during solidification. Small and spherical pores were therefore formed within the SLM-built samples. Besides, insufficient energy input also results in lack of fusion

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between adjacent tracks and successive layers [38], and consequently, unmelted particles [13] as well as crack-like pores [39]. Therefore, large and irregular pores were observed in the SLM-built sample at

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scanning speed of 1000 mm/s (Fig. 8c). The existing pores or voids play a negative role in the matrix continuity, which could induce stress concentration and subsequent crack nucleation under tensile stress. Therefore, the sample at high scanning speed presents relatively smaller elongation to failure. In contrary, the low scanning speed results in the improved ductility considering the highly densified microstructure of the as-built 316L stainless steel. Fig. 10 shows the SEM images of the melting pool at scanning speeds of 800 mm/s and 1000 mm/s. The microstructure in Fig. 10a and 10b indicates that the melting pool looks normal at scanning speed of 800 mm/s, while irregular melting pools with multiple boundaries were observed at scanning speed of 1000 mm/s. The observed melting pool boundaries at high scanning speed (Fig. 10b) increased the probability of brittle feature fracture, where crack nucleation and propagation occur. The improvement in ductility using a low scanning speed was attributed to the delay in crack initiation and

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propagation by decreasing the residual defects and melting pool boundaries.

Fig. 10 SEM images of the melting pool morphology at scanning speeds of 800 mm/s (a, c) and 1000 mm/s (b, d).

The zoomed-in melting pools shown in Fig. 10c and 10d indicate that the orientations of the dendrite/cellular structures varied at different locations within one melting pool, and were not exactly vertical to the melting pool boundary. This phenomenon can be explained by considering the coupling effect of the solidification condition, crystal structure, and convection flow inside the melting pool [26]. Highly-magnified SEM images of typical solidification structures are depicted in Fig. 11, where a distinct difference of the microstructure was observed. From the morphology of the surface vertical to the building direction (Fig 11a and 11b), the bright boundaries (cell walls) formed solidification cells with length and width of approximately 400 nm and 100 nm, respectively. The presence of

Journal Pre-proof elongated cells was more evident with high scanning speed of 1000 mm/s (Fig. 11b). It has been reported that dislocations and associated element segregations were concentrated at the cell walls [40]. One explanation for the formation mechanism of the solidification cell structure is the compositional fluctuations and constitutional supercooling theory during SLM [27]. Thus, it is thus reasonable to obtain smaller and more evident solidification cells by applying high scanning speed (Fig. 11b). The cross-sectional view (parallel to the building direction) in Fig. 11c and 11d show that spherical nanoparticles (indicated by red arrows) with size less than 100 nm were randomly located

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within the solidification cells. The formation of these nanoparticles was attributed to the high viscosity of the melted silicate, which tends to form a spherical shape during solidification and often

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remains undissolved in the structure owing to the low wetting tendency to steel [41]. The amount of

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nanoparticles in Fig. 11d was found to be higher than that in Fig. 11c, implying that the nanoparticle

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formation is related to the scanning speed. The nanoparticles were reported as a strengthening factor

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[16].

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for as-built 316L stainless steel [27], but their contribution was limited by the low volume fraction

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Fig. 11 Magnified SEM morphologies on the surfaces vertical (a, b) and parallel (c, d) to the building direction at scanning speeds of 800 mm/s (a, c) and 1000 mm/s (b, d). 3.3. EBSD analysis

It is well-known that the grain size and orientation play an essential role in controlling the strength of SLM-manufactured 316L stainless steel [42]. Fig. 12 shows the EBSD analysis of the samples at scanning speeds of 800 mm/s (Fig. 12a, c, and e) and 1000 mm/s (Fig. 12b, d, and f). Both coarse grains (~100 μm) and fine grains were found in the EBSD results shown in Fig. 12a and 12b. A continuous change (as indicated by color in Fig. 12) of grain orientations was also observed within the coarse grains, especially at high scanning speed (Fig. 12b), which reveals internal grain boundaries within the coarse grains. Based on the inverse pole figures in Fig. 12c and 12d, the grain growth did not exactly follow the most preferred growth orientation. This shift may be attributed to the 115° rotation of the scanning vector between successive layers, which led to the modification of

Journal Pre-proof the thermal gradient, and consequently the growth direction of grains. Fig. 12e and 12f show the misorientation angles between grains. The interfaces between grains where the misorientation angles ranged from 2°–10° were considered as low-angle grain boundaries (LAGBs, in red), and the interfaces with misorientation angles over 10° were defined as high-angle grain boundaries (HAGBs, in blue). The EBSD results shown in Fig. 12e and 12f indicate that a significant fraction of LAGBs appeared in the as-built sample, especially at high scanning speed of 1000 mm/s (Fig. 12f). It has been reported that the grain boundaries, especially those with high

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fraction of LAGBs, hinder the dislocation motion during tensile deformation, thus strengthening the

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manufactured samples [29,43].

Fig. 12 EBSD orientation maps (a, b), inverse pole figures (c, d), and grain boundary information (e,

Journal Pre-proof f) of the samples fabricated with the scanning speeds of 800 mm/s (a, c, and e) and 1000 mm/s (b, d, and f). Local misorientation angle from 0° to 2° associated with the LAGBs and HAGBs is shown in Fig. 13. The local misorientation angle was mapped pixel by pixel inside each grain, and a high fraction of misorientation angles in the range from 0.5° to 1.5° was observed (green and yellow ones shown in Fig. 13). For a single grain (indicated by the white rectangle in Fig. 13b), the variation of the local misorientation angle with respect to position is illustrated in Fig. 14. The local misorientation angle

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remained variable under 2°, and only a few points achieved the value of 2° (red line) and formed

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LAGBs (intragranular feature) and HAGBs (grain boundaries).

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Fig. 13 Local misorientation angles from 0° to 2° associated with LAGBs and HAGBs (black lines) at scanning speeds of (a) 800 mm/s and (b) 1000 mm/s.

Fig. 14 Misorientation angle variations with respect to position within one grain at scanning speed of 1000 mm/s. The LAGBs acted as a strengthening factor at the macro-level, and the large number of LAGBs

Journal Pre-proof impeded the movement of dislocations, thus increasing the strengthening effect through grain boundaries (HAGBs) [11]. As for the solidification cells, no local misorientation was reported across the majority of the solidification walls [43]. Instead, misorientation of up to 1.5° was observed inside the solidification cells. Therefore, the solidification cells worked as a strengthening factor that differs from LAGBs. Instead of serving to accommodate local misorientations as traditional dislocation walls [39], the solidification cells consisting of a high fraction of local misorientation provided a strong dislocation trapping and retention mechanism [43]. Accordingly, the solidification cells contributed to

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the high strength of the as-built samples by preventing dislocation motion at the micro-level. Based on the Hall–Petch strengthening relation with an assumed effective sub-grain length (relating to the size

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of solidification cells), the solidification cell structure is an important contributor to the high strength

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of the as-built 316L stainless steel.

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Aside from the local misorientations, the grain size distribution of the as-built samples was also studied, as shown in Fig. 15. Grains were defined using the HAGBs (>10°), and the grain size was

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defined as the average equivalent circle diameter of each grain area. The relatively fine grains were attributed to the high cooling rate while applying a high scanning speed. The relatively fine grains

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1000 mm/s.

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partially contributed to the high strength of the as-built samples obtained at high scanning speed of

Fig. 15 Grain size distribution of the as-built samples with scanning speeds of 800 mm/s (V1) and 1000 mm/s (V3). 4. Conclusions The present study examined the evolution of the mechanical performance of SLM-manufactured

Journal Pre-proof 316L stainless steel with the variation of process parameters, especially the scanning speed. The following conclusions could be drawn: (1) The UTS decreased slightly with increasing energy density, while the opposite tendency was observed for the elongation to failure. The maximum UTS of 707 MPa was achieved with scanning speed of 1000 mm/s, while maintaining good ductility (total elongation of 30%). (2) The scanning speed altered significantly the melting pool boundaries, residual pores, solidification cells, nano-inclusions, as well as grain size and distributions.

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(3) The maximum elongation to failure (55%) was achieved at low scanning speed of 800 mm/s, and

as well as the minimizing of melting pool boundaries.

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the strengthening of the high ductility was mainly attributed to the highly densified microstructure,

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(4) The cumulative effects of the fine cellular structure, high fraction of LAGBs, and refined grains

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contributed to the strengthening of the 316L stainless steel obtained at high scanning speed of 1000

Acknowledgments

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mm/s.

The authors acknowledge the financial support from the National Natural Science Foundation of

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China (NSFC, Grant No.: 51905306 and 51975335) and the China Postdoctoral Science Fund (2018M642650). Xiebin Wang is grateful for the financial support from the National Natural Science

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Foundation of China (NSFC, Grant No.: 51905310), Shandong Provincial Natural Science Foundation (Grant No.: ZR2018QEM001), Natural Science Foundation of Jiangsu Province (Grant No.: BK20180231).

CRediT authorship contribution statement Jiangwei Liu: Conceptualization, Validation, Formal analysis, Investigation, Writing - original draft, Writing - review & editing. Yanan Song: Investigation. Chaoyue Chen: Investigation. Xiebin Wang: Conceptualization, Methodology, Supervision. Hu Li: Writing - review & editing. Chang’an Zhou: Investigation. Jiang Wang: Conceptualization, Methodology. Kai Guo: Supervision, Project administration, Funding acquisition. Jie Sun: Project administration, Funding acquisition. Data availability The raw and processed data required to reproduce these results are available by contacting the

Journal Pre-proof authors.

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Journal Pre-proof Declaration of interests ☒ The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. ☐The authors declare the following financial interests/personal relationships which may be

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Graphical abstract

Journal Pre-proof Highlights: The scanning speed altered significantly the melting pool boundaries, residual pores, solidification cells, nano-inclusions, grain size and distributions. The high ultimate tensile strength was attributed mainly to the fine cellular structures and fine grains with low-angle grain boundaries.

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A low scanning speed results in an enhanced ductility due to the minimization of residual defects and melting pool boundaries.