Densification Behavior of 316L Stainless Steel Parts Fabricated by Selective Laser Melting by Variation in Laser Energy Density

Abstract

Selective laser melting (SLM) is an attractive manufacturing technique for the production of metal parts with complex geometries and high performance. This manufacturing process is characterized by highly localized laser energy inputs during short interaction times which significantly affect the densification process. In this present work, experimental investigation of fabricating 316L stainless steel parts by SLM process was conducted to determine the effect of different laser energy densities on the densification behavior and resultant microstructural development. It was found that using a low laser energy density below 50 J/mm³ produced an instable melt pool that resulted in the formation of unmelted particles, pores, cracks, and balling in the as-built parts with low densification. In contrast, the as-built parts at a high energy density above 200 J/mm³ showed irregular scan tracks with a number of pores and metal balls that decreased part density. The optimal laser energy density range was accordingly determined to be 58–200 J/mm³ by eliminating obvious SLM defects, which led to near full densification. The SLM samples fabricated using optimal parameters allowed observation of a microhardness of 280 Hv, ultimate strength of 570 MPa, and yield strength of 530 MPa that were higher than those of the as-cast and wrought 316L stainless steel.

1. Introduction

Additive manufacturing (AM), well known as rapid prototyping in the mid-1980's, is an advanced manufacturing technology to capable of making three-dimensional (3D) structure from a layer by layer materials deposition using digitally controlled machine tools^1–4). According to the ASTM standard⁵⁾, all type of materials can be processed by AM, polymers being the most studied to data. Recently, interest in AM processes for metal parts production is significantly increasing for applications in diverse industries including aerospace, aviation, automotive, and medical^4,6–10). One type of additive manufacturing technologies is selective laser melting (SLM). SLM allows the fabrication of 3D parts through the use of focused laser energy to powder beds directly from user defined 3D computer-aided design (CAD) models^4,11–13).

Generally, this technology is considered as one of the most promising manufacturing techniques for metallic materials due to its ability to produce complex geometries with unique design and high accuracy that can be used not only for the prototyping step but also for small series productions^2–4). During the SLM process, the metallic powders are completely molten by the laser beam, which enables the production of individual parts with high density and excellent mechanical properties^14–16). With respect to the material aspects, various investigations in SLM process have been conducted using 316L stainless steel (SS316L) powder because of its superior mechanical properties, including excellent corrosion and oxidation resistance, high heat resistance, and good weldability^17–20).

In the SLM process, it is well known that there are more than 100 parameters affecting the quality of final products²¹⁾. These parameters include not only size, composition, and flowability of initial powder particles but also laser spot size, scan speed, scan spacing, scan strategy, and layer thickness of the process conditions. In particular, the laser energy density, E (J/mm³) plays a decisive role in the development of the porosity and microstructure during the SLM process. In general, the laser energy density (E) can be calculated by the following equation²²⁾:   

\[E = \frac{P}{vht}\](1)

where P is the laser power (W), v is the scanning speed (mm/s), h is the hatch spacing (mm), and t is the layer thickness (mm). According to the eq. (1), the laser energy density E is simply controlled by the initial SLM process conditions of the laser powder, scanning speed, hatch spacing, and layer thickness. Thus, in this work, a parametric study was carried out to investigate the influence of the laser energy density on the densification and microstructural development of SS316L parts fabricated by the SLM process. The variation of the incident energy density was controlled by adjusting the laser scanning speed (v). The hardness and tensile properties of the fabricated samples were also evaluated, and consequently the suitable process window for SLM was properly determined.

2. Experimental Procedure

The powder materials used in this study was commercially available 316L stainless steel powder (CL20ES, Concept Laser GmbH, Germany). As seen in Fig. 1 (a) and (b), the powder particles reveal mostly spherical morphology while some irregular shaped particles, and contains a small number of satellites. Moreover, the particle size distribution is preferably narrow varying between 10 to 45 μm with an average diameter of 28 μm (see Fig. 1(c)). The SLM experiments were conducted using a commercial machine (Mlab-Cusing, Concept Laser GmbH, Germany) equipped with Yb:YAG fiber laser (λ = 1075 nm) with an effective power of 100 W.

Fig. 1

(a) SEM micrographs of the as-received SS316L powders and (b) their cross section, and (c) particle size distribution.

The SLM process is carried out in two ways as presented in Fig. 2. Small SS316L cubes (Fig. 2(a)) with size of 10 mm × 10 mm × 10 mm were at first prepared using different laser energy densities ranging from 10 to 1000 J/mm³ induced by variation of laser scanning speeds up to 3000 mm/s, in order to identify the suitable process parameters compared with the porosity and microstructure of as-built parts. Next, to analyze the mechanical properties, the subsize tensile specimens were prepared following the ASTM E8²³⁾ procedure from the thin wall structure (Fig. 2(b)). The specimens were fabricated using selective processing parameters that produced the highest density: 2.5 mm in thickness, with a gauge length of 25 mm and a gauge width of 6 mm. All the specimens were prepared along with a Z-increment (vertical) at a constant laser powder of 90 W, a focused laser beam diameter of 110 μm, hatch spacing of 80 μm, and layer thickness of 25 μm. During the SLM process in the present work, a 316L stainless steel base plate was used as a substrate for the deposits, and the line scanning strategy that laser moved bidirectionally (zigzag) across the surface was applied as schematically illustrated in Ref. 24). The build chamber was filled with an inert Ar atmosphere (Ar, 99.999%) to maintain the oxygen content less than 0.4%. Within the study, the term “as-built” refers to the SS316L parts prepared by the SLM process without any post treatments.

Fig. 2

As-built SS316L specimens in this study: (a) cubic and (b) thin-wall specimen.

The phase analysis of raw materials and as-built SS316L samples was carried out by x-ray diffraction (XRD, AXS Bruker D5005 diffractometer, Germany) using Cu Kα target (λ = 1.5418 Å) in the angular range (2θ) 20° to 80° at a scanning peed of 4°/min. In order to evaluate the compositional change during the SLM process, the SS316L samples before and after SLM were chemically analyzed using ICP (Spectro Arcos ICP-AES, Germany), C/S (Eltra CS-800, Germany) and N/O (Eltra ON900, Germany) measurements. The relative density of as-built samples was measured by Archimedes' method using distilled water. Microstructural observation was performed by means of optical microscope (OM, Nikon Microphot Fax, Japan) and field emission scanning electron microscope (FE-SEM, Tescan MIRA-3, Czech Republic) equipped with an electron backscatter diffraction (EBSD) system. To reveal the microstructure, the as-built samples were subjected to mechanical polishing and chemical etching in diluted aqua regia solution (HCl:HNO₃:H₂O = 3:1:4) for 40 s, and particularly for EBSD imaging the samples were electro polished using a 5% perchloric acid solution. Tensile tests were performed in a universal testing machine (RB 301 UNITECH-T, R&B Inc., Republic of Korea) at room temperature with a head speed of 1 mm/s, and the Vickers hardness was measured on the cross sections (XY plane) using a Mitutoyo HM-211 microhardness tester (Mitutoyo Co., Japan) at a loading of 0.5 kg and an indentation time of 20 s. Each mechanical test was performed at least three times under the same condition, and the average value was calculated. The tensile fracture surfaces were also estimated by SEM.

3. Results and Discussion

Figure 3 shows the effect of the applied laser energy on the relative density of the as-built samples while maintaining the constant value of laser powder, layer thickness, and hatching spacing of 90 W, 25 μm, and 80 μm, respectively. As shown in Fig. 3, the overall trend of increasing density from 83 to 99% of the theoretical density (% TD) was observed with increasing E up to about 58 J/mm³, after which the density was maintained at near full density of > 98.5% TD. This is due to the relatively high E caused by lower scanning speed that promotes a stable melt pool by increasing molten materials temperature. This may also lower the surface tension and enhance the wetting characteristics of molten materials resulting in smooth scan tracks and dense structure^22,25). However, when the E was applied over 200 J/mm³, the density of the as-built samples decreased again. The reason for this phenomenon is that the SLM process at very high energy, E > 200 J/mm³ in this work, may increase the risk of balling and dross formation in the melt pool, thereby lowering the density and surface quality^3,26). In this context, there is a critical E value that follows the proportional increase in the part density with increasing laser energy density. Therefore, the proper E range can be determined to be about 58–200 J/mm³, which allows the high density above 98.5% TD to be obtained.

Fig. 3

The variation of relative density of the as-built cubic samples with increasing laser energy density.

Phase identification of the raw powder as well as SLM processed sample with the highest density were conducted as shown in Fig. 4, particularly the analysis plane of SLM sample is perpendicular to the building direction (Z-axis), i.e., XY plane in Fig. 2(a). It can be seen that all XRD patterns exhibited peaks corresponding to the fcc austenite phase (γ-Fe, JCPDS-ICDD, PDF No. 33-0397) without any detectable presence of ferrite. In addition, it was found that the diffraction peaks after the SLM process were slightly broader compared to that of the raw materials, which suggested the presence of residual stresses and dislocations introduced during the SLM process. It is a general feature in laser-processed materials because it contains rapid heating/cooling stages during the process, which results in an increase in the residual stress and dislocation density as previously reported^27,28).

Fig. 4

XRD patterns of the SS316L raw powder and SLM sample produced at a laser energy density of 58.4 J/mm³.

In the SLM process, laser energy input often causes an increase in raw materials vaporization, spatter generation, and chemical contaminations, which directly affects the composition and quality of the as-built samples^3,26). As a consequence, a chemical analysis of SS316L before and after SLM was conducted, as shown in Table 1. It is clear that the overall composition of SS316L did not substantially change during the SLM process. Also, the elements of C and O for the as-built samples were similar in chemical composition to that of the raw materials, which implies that no serious contamination occurred during the process in the Ar atmosphere.

Table 1 Chemical composition of the SS316L raw powder and the as-built sample at a laser energy density of 58.4 J/mm³ (mass%).

	Element
Samples	Fe	Cr	Ni	Mo	Si	Mn	S	P	C	O
Powder	Bal.	18.13	12.98	2.36	0.66	1.69	0.005	0.010	0.02	0.045
SLM, E = 58.4 J/mm3	Bal.	17.52	12.11	2.31	0.62	1.24	0.004	0.012	0.01	0.044

To evaluate the microstructural development during SLM process, the comparable microstructures of the top surface for the as-built cubic samples at different laser energy densities are shown in Fig. 5. These showed in good agreement with the results of Fig. 3. At a very low energy density (E = 29.2 J/mm³) induced by a high laser scanning speed of 1600 mm/s, discontinuous laser tracks were clearly visible, and large and highly irregular pores were located at the melt pool boundaries, as depicted in Fig. 5(a). In addition, unmelted and loosely bonded particles were observed within an elongated pore (Fig. 5(a), ×500). This can be mainly attributed to the lower penetration of laser energy, which leads to the reduced melt pool size and increases the unmelted powder particles, as there is not enough to ensure sufficient bonding between the scan tracks and/or layers²²⁾. The highest part density (99.2% TD in Fig. 3) can be attained as an increase in laser energy density from 29.2 to 58.4 J/mm³. At this point, most of the defects such as imperfect bonding between the scan tracks, large and irregular pores disappeared. Instead, only a few microvoids were detectable in the microstructure as demonstrated in Fig. 5(b). Normally, the small pores occurring in the SLM microstructure were reported as the gas porosity and solidification shrinkage⁹⁾. On an increasing the E up to 116.9 J/mm³ as indicated in Fig. 5(c), small and spherical pores generally surrounded by dense regions were observed over the sample. Here, the part density was measured to be about 99% TD achieving near full density, but some defects as microcracks and voids were still remained in the microstructure. Figure 5(d) shows the microstructure of the SLM sample fabricated at very high E (> 200 J/mm³). As observed in this figure, it was found that numerous pores and cracks were detected again inside the melt pools, and especially several closed and large pores up to 20 μm were present. This excessive energy input could result in a large molten pool with high liquid phase content affecting the formation of the elongated and more irregular melt pool, which causes the quality of final products to deteriorate by formation of large pores^29,30).

Fig. 5

OM micrographs showing the microstructure of the as-built samples at different laser energy densities of (a) 29.2, (b) 58.4, (c) 116.9 and (d) 233.8 J/mm³.

Melt flow behavior is another factor that may have the tendency to influence the porosity development during SLM process and further affect the surface quality^31,32). Figure 6 elucidates the representative top surface morphologies of SLM cubic samples manufactured using different laser energy densities of 29.2, 116.9, and 233.8 J/mm³, respectively, which contains the laser scanned tracks indicating the traces of melt flow during the SLM process. As seen in the figures, the SLM samples produced under the three conditions featured long bidirectional scanning tracks and the track width was confirmed to be approximately 85 μm, which is well in agreement with the hatch spacing of 80 μm, but the surface morphology was different. At very low E level of 29.2 J/mm³, discontinuous and spacing were observed in between the scan tracks, and a few pores were found in the irregular-shaped melt flow traces, as shown in Fig. 6(a). Furthermore, a large amount of unmelted (of partially melted) powder particles and metallic balls were clearly visible throughout the sample. This can be explained by the insufficient melting of the powder materials under the low energy input that generated the instable melt flow which was solidified before adequate metallurgical bonding with the previous layer, incurring the defects formation²²⁾.

Fig. 6

SEM micrographs of top surface of the as-built SS316L samples at different laser energy densities: (a) 29.2, (b) 116.9 and (c) 233.8 J/mm³.

As the E was settled at 116.9 J/mm³ (Fig. 6(b)), the surface showed generally regular laser scan tracks and an overlapped melt pool region with a relatively smooth surface, but still contained a small amount of pores and spherical balls. According to the research from Gu and Shen^33,34), high laser input during the SLM process increases the molten materials temperature that establishes a stable melt pool with favorable melt flow characteristics including surface tension and wetting ability. This process results in sound metallurgical bonding and smooth melt tracks. At a laser energy density of 233.8 J/mm³ shown in Fig. 6(c), the laser scan tracks were continuous and overlapped in a well condition but a number of metallic balls were observed on the surface as compared to that of Fig. 6(b). This high energy input causes the longer dwelling time of laser beam on the surface of the molten materials that results in an increase in the temperature of the melt pool that accordingly decreases the dynamic viscosity of the liquid phases in the melt pool. These combination effects lead to the overheating of liquid, and relatively unstable melt pool development. Under this condition, a number of small-sized liquid droplets tend to splash from the melt pool front being solidified, which leads to material spatters and balls forming on the surface^33,34). Sometimes there is a contradictory relationship between laser energy input and surface quality as elucidated in Fig. 6(b) and (c). Therefore, finding the optimal process parameters, i.e. laser energy density in this study, is very important in SLM process. Based on the above results, it is reasonable to conclude that the SLM process at a very low (E < 50 J/mm³) and high (E > 200 J/mm³) energy density is unsuitable for part building because it leads to pore generation and balling and dross formation that directly affects the quality of the final product.

EBSD was further performed to evaluate the microstructural characterization, especially the size and distribution of the grain structure, of the SS316L parts produced by SLM. Figure 7 discloses the representative EBSD orientation map of the indexed area for as-built samples at different E of 58.4 and 116.9 J/mm³, respectively. Both of the samples showed high density above 99% TD and relatively uniform microstructure, as appeared in Fig. 3 and Fig. 5. However, a significant difference in the grain size between the samples was clearly observed in Fig. 7; the average size was measured to be about 30 μm in the as-built sample for relatively low E (Fig. 7(a)), whereas the as-built sample under high E was composed of coarsened grains above 100 μm (Fig. 7(b)). This may have resulted from the different solidification and cooling rate of a melt pool in the SLM process, which are largely dependent on the thermal gradient in the laser scan direction^35–37).

Fig. 7

EBSD mapping of the as-built SS316L samples under different E of (a) 58.4 and (b) 116.9 J/mm³.

As reported by Li and Gu³⁸⁾, higher laser energy yields a high temperature and long life time of the molten pool, which gives rise to a large amount of liquid formation, and therefore it can be expected to produce smaller thermal gradient and slower cooling rate. As a consequent, this thermal behavior under high E could be detrimental to the refinement of microstructure and resultant mechanical properties of the final component. Furthermore, at low E, the insufficient melt laser energy input to the powder layers could lead to a small amount of liquid formation and large thermal gradient induced by high cooling rate causing defect formation in the as-built parts, which corresponded to the E below 50 J/mm³ in this work, as described above in Fig. 3 and Fig. 5(a). Thus, it is worth noting again that the determination of the proper level of laser energy density is a critical factor for the SLM process in order to acquire the final component with high density, fine microstructure and resultant mechanical properties at the same time. The in-depth aspect of the microstructural evolution and related mechanical properties in the SLM process is currently under investigation for the SS316L and will be reported in a follow-up study.

Figure 8 depicts the correlation between hardness values and applied laser energy density of the as-built cubic samples, in which a clear tendency is observed where the hardness increases with part density shown in Fig. 3. The individual points of the hardness value are indicated an average of ten measurements with the standard deviation (except the maximum and minimum values). As seen in this figure, there was a general increase in the hardness value with increasing energy density, and a maximum value of 240 Hv could be achieved at 58.4 J/mm³. Additionally, the average hardness value of the dense samples (> 98.5% TD) was measured to be about 235 Hv with a standard deviation of 3%, which corresponded to the typical hardness values from the previous experimental results of Ref. 39). A further increase in E, however, exhibited a considerable decrease in hardness to 187 Hv related to the porosity changes in the as-built samples induced by the laser energy input where the standard deviation increased to > 10%. It is worth noting here that changes in the standard deviation of the hardness mostly arise from the structural homogeneity, and synchronously have an influence on the mechanical properties of the final products^40,41). The as-built SS316L samples at both very low and high E were, therefore, believed to compose a relatively less uniform microstructure than that of the samples produced at the proper E condition.

Fig. 8

Microhardness value of the as-built SS316L part with different energy density.

Figure 9 illustrates the stress-stain curves for the as-built SS316L specimen with the highest density corresponding to a laser energy density of 58.4 J/mm³ (refer to Fig. 3). This typically provides the mechanical behavior including elastic deformation, yielding followed by plastic deformation with work hardening, and final fracture. The testing results of SS316L samples produced by SLM and conventional methods are summarized in Table 2. The as-built specimen showed higher yield strength (0.2% offset yield strength, σ_0.2) and ultimate tensile strength (UTS) compared to those achieved from other manufacturing processes. The as-built samples also satisfied the literature values in the UTS range between 525 and 623 MPa for the cold finished wrought SS316L⁴³⁾. On the other hand, the elongation was lower than that of the as-cast sample, but still attained a relatively high elongation of more than 40% even in the as-built condition. This is mainly attributed to the good metallurgical bonding between the adjacent layers and melt tracks that is well developed within the selected laser energy density in this work. The high ductility can also be explained by the stable austenitic microstructure of SS316L not undergoing a phase transformation after the melting and subsequent solidification during SLM process⁴²⁾, as described in Fig. 4.

Fig. 9

Typical stress-strain curves of the as-built SS316L specimen at a laser energy density of 58.4 J/mm³.

Table 2 Comparison the tensile properties for SS316L manufactured by SLM and by conventional processing methods⁴³⁾.

Samples	σ0.2 (MPa)	UTS (MPa)	Elongation (%)	References
SLM, E = 58.4 J/mm3	531.8 ± 2.7	573.3 ± 5.8	41.3 ± 1.4	this study
As-cast	262	552	55	43)
Wrought (cold finished)	255	525–623	30	43)
Wrought (hot finished)	170	480	40	43)

It should be noted here that the considerable increase in yield strength of 530 MPa can be achieved at SLM processed SS316L in the as-built condition whereas the typical yield strength of SS316L processed by conventionally manufacturing techniques is less than 300 MPa. This improved property seems to be mainly caused by formation of the fine microstructure, as has been reported in various works^44–46). Generally, grain refinement leads to an increase in both the yield strength and tensile strength. The yield strength of the materials is more related with the grain size in accordance with the Hall-Petch relationship⁴⁷⁾ where the smaller grain size gives higher material yield strength. It is also well known that the interaction of high density dislocations can improve the yield strength of the material. The SLM process, in this study, involves high cooling rate and rapid solidification that possibly produce the fine structures as shown in Fig. 7(a). This may subsequently lead to the generation of large residual stresses and high dislocation densities⁴²⁾, which in turns improves the mechanical properties of the resultant SS316L parts, as discussed above in Fig. 4.

Figure 10 presents the fracture surface of SLM SS316L samples produced at E of 58.4 J/mm³ after tensile tests, as shown in Fig. 9. It can be seen in Fig. 10(a) that the sample shows ductile fracture morphologies without large voids and cracks that are considered critical sites for initiating the material failure during the tensile test. A higher magnification view of Fig. 10(b) displays dimple structures with uniform size distributed on the fracture surface, which confirms the main fracture mechanism is ductile fracture.

Fig. 10

SEM micrographs of tensile fracture surfaces of the as-built SS316L specimen corresponding to Fig. 9.

At low energy inputs during the SLM process, unmelted or partially melted particles are occasionally preserved in the boundary of the adjacent melt pools due to the lack of heat transfer, resulting in the formation of large and irregular shaped voids, as displayed in Fig. 5(a). On the contrary, when the laser energy density is excessive, the undesirable metallurgical pores are generated by entrapment of gases or improper closure of a keyhole⁴⁸⁾, as demonstrated in Fig. 5(d). The presence of these microstructure defects weakens the bonding between the consecutive melt tracks or neighboring layers, which have a significant effect on the mechanical performance⁴⁹⁾. In this study, the SLM defects can therefore be effectively eliminated by applying the optimal laser energy density of 58.4 J/mm³, improving the mechanical properties. Moreover, the selected laser energy density of 58.4 J/mm³ involves a grain refinement effect that enhances the mechanical properties of as-built SS316L parts because of the increased grain boundaries acting as barriers for dislocation motion⁵⁰⁾.

4. Conclusions

In this study, 316L stainless steel parts were successfully fabricated by applying a selective laser melting process with different laser energy densities. The relationship between the applied laser energy, densification, and microstructure of SS316L parts processed by SLM was effectively established. The results are summarized as follows:

(1) The densification behavior of SLM processed SS316L parts was effectively controlled by the applied laser energy density (E). Nearly fully dense (above 98.5% of the theoretical density) samples were obtained at E ranging from 58 to 200 J/mm³ by eliminating the formation of macro defects. Furthermore, the as-built part with the maximum density of 99.2% TD was produced at E of 58.4 J/mm³ without any significant chemical changes.

(2) Several SLM defects such as pores, cracks, and irregular melt tracks were found in the microstructure of as-built SS316L samples at both relatively low and high energy density region, E < 50 and E > 200 J/mm³, respectively. The microstructure of the as-built samples at 58.4 and 116.9 J/mm³ revealed sound metallurgical bonding and smooth melt tracks due to the formation of a sufficient melt pool, but a finer microstructure could be obtained at energy density of 58.4 J/mm³

(3) In this study, the as-built SS316L parts with a microhardness of 240 Hv, UTS of 570 MPa, EL of 40%, and σ_0.2 of 530 MPa were achieved using the laser energy density of 58.4 J/mm³. In particular, the yield strength was considerably higher than that of conventionally manufactured samples. Consequently, the SLM process in this work is a reliable process to fabricate SS316L parts that can result in the inhibition of the SLM defects as well as development of the grain refinement, thereby improving the mechanical properties.

Acknowledgements

This study was supported financially by Fundamental Research Program of the Korean Institute of Materials Science (KIMS). This work was also supported by the National Research Council of Science & Technology (NST) grant by the Korea government (MSIP), (No. CRC-15-03-KIMM).

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