Introduction

Directed energy deposition (DED) is a class of additive manufacturing (AM) processes that allow the production of metallic components by melting, by means of a focused energy source, the metallic material which is directly fed in specific areas of the building volume. DED technologies gained a large interest mainly thanks to the possibility to repair metallic components and produce large and functionally graded parts. In particular, laser powder-DED (LP-DED) uses a laser beam as energy source and the material in the form of powder. LP-DED was recently used to process various alloys such as AlSi10Mg, Ti-6Al-4V, Inconel 625 and AISI 316L stainless steel (SS) (Ref 1,2,3,4). The results showed that it is possible to produce dense and crack-free parts with these powders and that, thanks to the rapid cooling, the as-built components are constituted by a fine microstructure and interesting mechanical properties.

In particular, the microstructure and mechanical properties of LP-DED 316L parts have been investigated by several authors in recent years (Ref 5,6,7) The main findings are that the as-built microstructure is strongly correlated with the thermal history to which the material is subjected during the AM process and in particular with the high heating and cooling rates as well as the remarkable thermal gradient. These phenomena cause the solidification of a distinct microstructure made of large grains containing fine cellular dendrites with a primary cellular arm spacing (PCAS) of a few microns (Ref 5). The size and the morphology of these dendrites strongly vary within a sample and are mainly connected to the location within the melt pool. This complex dendritic structure is characterized by the presence of two main phases: the face-centered cubic (FCC) austenite (γ) phase and body-centered cubic (BCC) ferrite (δ) phase. The presence of this unique microstructure explains the high mechanical properties of as-built 316L LP-DED parts. The high properties are indeed mainly due to the reduced size of dendrites, the presence of the hard δ-ferrite phase and the presence of a dense dislocation network within the cells (Ref 8, 9).

Notwithstanding this, LP-DED components are characterized by the presence of residual stresses which are also due to the peculiar thermal history to which the material is subjected while being processed. The presence of such residual stresses affects important characteristics of the produced parts such as tensile and fatigue properties; thus, the integrity and the lifetime of components are influenced. Wagner (Ref 10) showed that the presence of tensile residual stresses reduced the fatigue live due to the high driving force to cracks propagation. Vrancken et al. (Ref 11) using finite element model (FEM) observed that the presence of residual stresses was the mainly factor that affects the anisotropic behavior of the components. Withers and Bhadeshia (Ref 12) demonstrated that residual stresses are mainly caused by the thermal misfit between two adjacent regions and by the dynamic nature of thermal phenomena in laser-based processes. According to Mercelis and Kruth (Ref 13), two mechanisms are primarily responsible for the generation of residual stresses: the temperature gradient mechanism and the cooling down phase. The temperature gradient mechanism is caused by the high-temperature gradient near to the laser spot, whereas in the cooling down phase model the stresses are generated by material shrinkage during the solidification or cooling. However, this contraction is limited by the underneath solid material.

In recent years, several studies have been carried out to evaluate the residual stress in LP-DED samples. Rangaswamy et al. (Ref 14) measured the residual stresses of two LP-DED stainless steel samples using the neutron diffraction method. The samples were characterized by different geometries, i.e., a thin wall and a pillar of square cross section. They demonstrated that a compression state was observed at the center of both samples, whereas a tension state was obtained near the edge. Later, Rangaswamy et al. (Ref 15) measured the residual stress distribution of two samples with a rectangular and square cross section using neutron diffraction and the contour method. They observed that the stresses were almost uniaxial along the building direction and that they were almost independent from deposition direction. Woo et al. (Ref 16) studied the residual stress distribution obtained in 316L SS-P21 functionally graded material samples produced by LP-DED varying the deposition strategy. They showed that the island deposition strategy significantly reduced the amount of residual stresses, but a high level of lack of fusion defects was observed. The presence of such defects is due to poor applicability of the island strategy to the LP-DED process due to the necessity to have some overlap between neighboring islands. Moreover, Saboori et al. (Ref 17) showed that residual stresses of 316L LP-DED samples were characterized by an oscillatory nature and that higher value of stresses was obtained on the lateral surface.

Residual stresses developed in LP-DED and generally in AM samples and components during the building process can be reduced through post-processing heat treatments thanks to the initiation of diffusional and relaxation phenomena. These phenomena might also cause the modification of the microstructure and, in some cases, the reduction in the tensile properties. The thermal treatments conditions for LP-DED 316L stainless steel are, however, not yet defined. It is therefore fundamental to establish the most convenient heat treatment that allows the residual stress relaxation and maintains the correct microstructural features as much as possible, so that the mechanical properties would still be high. Up to now, only a few authors investigated the effect of heat treatments on the mechanical performance and microstructure of LP-DED 316L steel. Yadollahi et al. (Ref 6), in particular, performed a heat treatment for 2 h at 1150 °C followed by air cooling on 316L LP-DED samples and obtained a recrystallized microstructure constituted by equiaxed grains. Heat-treated (HT) samples were, however, characterized by about 30% lower Vickers hardness and yield strength (YS). The authors attributed the reduction in mechanical properties to the grain growth and to the reduction in δ-ferrite content. Morrow et al. (Ref 18) heat-treated LP-DED 316L tensile bars at 1060 °C for 1 h and also observed a reduction in YS and ultimate tensile strength (UTS), but higher elongation (ε) values. Up to now, to the best of the authors' knowledge, lower temperature post-processing heat treatments, generally used for laser powder bed fusion parts, have not been studied for LP-DED materials.

In the present work, the effect of distinct deposition strategies on residual stress trends and mechanical properties of 316L SS LP-DED as-built samples was investigated. The differences in residual stresses and hardness along the sample heights were evaluated by performing the measurements at different distances from the substrate. Furthermore, the effect of stress relieving heat treatments on the microstructure, mechanical properties and residual stress of LP-DED 316L samples is presented and discussed. Finally, in order to assess the effect of the heat treatments on the tensile properties, tensile tests were carried out on as-built and heat-treated samples.

Materials and Methods

Material

A fresh commercial 316L SS gas atomized powder with a particles size in the range 50-150 µm was used for the production of the samples. Further information about the powder used in this work can be found elsewhere (Ref 19). Scanning electron microscope (SEM) micrographs of the as-received powder are reported in Fig. 1 and show that the particles have a spherical morphology, but present some satellites.

Fig. 1
figure 1

SEM micrograph of AISI 316L gas atomized powder

LP-DED Process and Heat Treatments

The samples were built with a prototype LP-DED system equipped with an Yb laser (YLS 3000, IPG Laser) on a three-axis CNC unit to control the movement of the XY deposition table. A multi-nozzle (4 ways) is used to deliver the powder into the melt pool. All the samples were performed at fixed parameters of laser power, laser spot diameter, hatching distance, travel speed and layer thickness which lead to porosity values lower than 0.3%. Two types of samples were built for these experimental activities, namely cubes, having 20 mm side, and bars, having the following dimensions 12 × 12 × 93 mm3 (Fig. 2a and b). A total of 12 cubes and 6 bars were built. Cubes were used for microstructural and stress analyses, and bars instead were used for tensile tests. Two deposition strategies, characterized by a different rotation angle between the layers (60° and 90°), were used for the deposition process (Fig. 3a and b). Finally, some samples underwent heat treatments at 600 and 800 °C for 2 h followed by air cooling. These temperatures were selected in order to avoid the recrystallization of the microstructure and the consequent reduction in mechanical properties.

Fig. 2
figure 2

Schematic representations of (a) cube and (b) bar samples

Fig. 3
figure 3

Schematic representations of the (a) 0°-60° and (b) 0°-90° rotated deposition strategy

The samples were identified using a three-field code based on the following criteria:

  • The first field represents the shape of the sample; in this work, cubic and bars samples were produced, and thus, the shape was identified by the letters C and B;

  • The second field identified the deposition strategy, and the numbers 0060 and 0090 identified the 0°-60° and the 0°-90° deposition strategy, respectively;

  • The third field identified the heat treatment adopted; AB means that the sample was analyzed in the as-built condition, and 600 and 800 represent the temperatures used in the heat treatment.

Table 1 summarizes the identification of samples and the process parameters used for the production.

Table 1 AISI 316L sample identification codes

Microstructural Characterizations

The samples were wire electrical discharge machining (WEDM) cut along the Z direction into two parts, as represented in Fig. 2(a). Part I was used for microstructure and hardness analysis, whereas part II was used for residual stress evaluation. Previous studies showed that wire EDM method is recognized as a suitable process for sample cutting due to the very small area influenced and modified by the cutting process (Ref 20, 21). After cutting operation, half of each sample was ground with 600-1200-2000 and 4000 grit SiC paper and polished using 3 and 1 µm diamond pastes. The polished surfaces were then observed by means of a Leica DMI 5000 optical microscope (Leica microsystems, Germany). The mean porosity value was evaluated by means of image analyses on 20 images taken at 100 × by means of the ImageJ (National Institutes of Health, USA) software.

The samples were then etched with a Kalling’s n2 solution for 10 s, and the microstructure was observed using the optical and the Phenom XL SEM microscope (Thermo Fisher Scientific, USA). X-ray diffraction (XRD) measurements were taken on the XY plane of the block samples using an X'pert Philips diffractometer (Cu Kα) in a Bragg–Brentano configuration with 2θ range = 20° − 110° (operated at 40 kV and 40 mA with a step size 0.013 and 35 s per step).

Residual Stress Measurements

The residual stresses were evaluated using the incremental hole-drilling strain-gauge method (IHDM). After the WEDM cut operation, the stress of the material is partially relaxed. However, as demonstrated by Bueckner (Ref 22), the change in residual stress distributions after a cut is determined by the components of stress normal to the cut plane. Later, Pagliaro et al. (Ref 23) showed using a finite element model that the global stresses were altered only in a region near the cut. In IHDM, only the stresses parallel to the cutting plane are evaluated. Therefore, the residual stresses measured by IHDM were not altered (Ref 24). Schajer (Ref 25) demonstrated that IHDM is recognized as one of the most efficient ones for evaluating residual stress distribution, in terms of cost, accuracy and versatility. In IHDM, a drilling cutter is used to produce a hole into the material to be tested. The hole is produced through a sequence of drilling steps. At each drilling step, a certain quantity of stressed material is removed. Consequently, a redistribution of residual stresses in the material around the hole is obtained and this causes a localized deformation. This deformation is acquired at each drilling step by at least three strain gauges, and these deformations are used to back-calculate the residual stress values in compliance with the procedure detailed in the E837-13a ASTM Standard (Ref 26). The volume of the material removed from the sample during the test is almost negligible; consequently, the IHDM is considered a semi-destructive method.

In this work, a RESTAN-MTS3000 (SINT Technology S.r.l, Italy) incremental hole-drilling system was used. A 1.8-mm-diameter coated carbide end mill with an inverted cone shape was adopted to drill a 1.2-mm-depth flat-bottomed hole. The relaxed deformations were acquired using a K-RY61-1.5/120R (HBM Italia S.r.l., Italy) Type B three-element rosette connected to a specialized amplifier. The surface of the specimen was manually prepared for the installation of the strain-gauge rosette using 200 and 400 grit SiC paper and wiped from contaminants and dust. The rosette was then bonded using Z70 cold curing superglue with the BCY01 (HBM Italia s.r.l., Italy) accelerator. The rosette cables were then blocked using an X60 (HBM Italia s.r.l., Italy) bicomponent cold curing glue. Figure 4(a) shows a schematic representation of the arrangement of strain-gauge rosettes on the specimen. Then, the sample was glued on a fixed platform and successively the end mill was aligned in the center of the strain-gauge rosette. In Fig. 5(a) and (b), a sample prior and during the measurement phase is illustrated. Incremental hole-drilling tests were carried out by executing a sequence of 24 steps each with a penetration depth of 50 µm; the relaxed deformation was acquired after each increment. The result of the drilling phase is visible in Fig. 5(c). The automatic RSM software (SINT Technology S.r.l, Italy) was used to acquire the deformation for each increment. The acquired deformations were then introduced into the EVAL software (SINT Technology S.r.l, Italy) in order to compute the residual stress profiles. The back-calculation was carried out in order to evaluate the stresses along specific directions. The selected directions (Fig. 4) correspond in detail to the building direction (i.e., z-axis), and the normal to the plane in which the rosette was applied (i.e., x-axis). As the selected directions are not necessarily principal directions, the outputs of the stress computation consist of two normal stresses (acting parallel to the two selected directions, respectively) and a shear stress τxz (acting on the plane defined by the two selected directions).

Fig. 4
figure 4

Schematic representation the arrangement of strain-gauge rosettes on part II of the sample

Fig. 5
figure 5

(a) Samples after the installation of the rosette (b) sample during measurement phase and (c) result of the drilling phase

Mechanical Testing

Vickers microhardness values were evaluated on the XZ cross section by means of a Leica VMHT indenter (UHL Technische Mikroskopie GmbH, Germany). The hardness trend along the Z direction was evaluated by performing the indentations with 300 g and 15 s. Indentations were performed in the sample’s center every 1 mm in order to evaluate the hardness mean values and standard deviations.

Finally, tensile samples were tested using a Zwick Z100 tensile machine using 8 × 10−3 s−1 as strain rate. The samples were machined based on the ASTM-E8 standard (Ref 27) from the AB and HT bars as represented in Fig. 2(b). Four samples per condition were tested.

Results and Discussion

In the following paragraphs, the results obtained on the AB and HT samples are described. Firstly, the results in terms of microstructure and the hardness are presented. Then, the residual stress distribution with respect to the depth is described.

As-Built

Porosity of C-0090-AB and C-0060-AB samples was measured to be 0.28 ± 0.11 and 0.23 ± 0.08%, respectively, showing that in both cases high density values could be achieved and that the deposition strategy does not have a significant effect on the consolidation of these samples. The microstructure of both samples is made of large grains containing austenite (γ) dendrites surrounded by the metastable δ-ferrite phase. No clear differences in terms of microstructure between samples built with different deposition strategies were observed, and the typical microstructure of the XZ cross section of AB samples is reported in Fig. 6. Very fine dark dots were observed in both samples (Fig. 6c and d), and the presence of similar features in AM 316L samples was revealed in previous literature works (Ref 19, 28, 29). Most of the experimental activities are revealed by means of electron-dispersive spectroscopy (EDS) and transmission electron microscopy (TEM) that these inclusions are oxides. Saboori et al. (Ref 19) for example, mentioned the presence of Mn-Si-rich oxides in LP-DED samples, and Eo et al. (Ref 28) reported the formation of Mn-Si-Cr oxides.

Fig. 6
figure 6

SEM micrograph of a AB 316L LP-DED built with (a), (c) 0°-90° and (b), (d) 0°-60° deposition strategy cross section

The hardness values were evaluated on the XZ cross section on two samples built with different deposition strategies. The hardness trend, reported in Fig. 7, shows that higher hardness values were measured in the lower part of the samples. This effect was previously observed by Javidani et al. (Ref 1) in AlSi10Mg LP-DED samples, and it is mainly due to the different melt pool cooling rates at distinct heights. The lower part of the samples has, in fact, a finer microstructure due to a high cooling rate, while the upper part is characterized by larger dendrites due to the lower cooling rate. The hardness mean values of C-0090-AB and C-0060-AB samples were very similar and measured to be 239.7 ± 7.3 Hv and 234.4 ± 8.0 Hv, respectively, confirming that the deposition strategy does not strongly affect the microstructure and the hardness values.

Fig. 7
figure 7

Hardness trend along the Z direction of AB samples

The most representative stress–strain curve of AB 316L samples is reported in Fig. 8. The mean YS and UTS values are 522 ± 7 and 662 ± 5 MPa. The elongation is 37 ± 1%. The high YS and UTS values of 316L AB samples are in line with literature data and are mainly due to the unique microstructure. As demonstrated by previous works, the high tensile properties are due to the fine microstructure, to the presence of the hard high-temperature δ-ferrite phase and by the high dislocation network of the as-built material (Ref 7, 30). The dense dislocation network explains also the poor strain hardening of the alloy (Ref 31).

Fig. 8
figure 8

Tensile curve of 316L LP-DED samples in the AB condition

The results of the numerical elaborations of the HDM measurements, taken according to the previously described procedure, are presented hereafter. The depth profiles of the normal residual stresses and of the shear stress (along the y direction) in the investigated samples are illustrated as a function of the deposition strategy and built height (distance from the substrate).

Figure 9 shows the comparison between residual stress along x-axis (σx) and along z-axis (σz) for AB samples produced using two different deposition strategies measured at z = 16 mm position. As visible in Fig. 9(a), C-0090-AB sample showed an initial value of σx of about 230 MPa; then, for a depth between 0.1 mm and 0.8 mm, a small oscillation of stress value around the zero value was observed. For a depth between 0.8 and 1 mm, the stress value abruptly increased and reached the value of 160 MPa. On the other hand, in the C-0060-AB sample, the amplitude of the oscillation of σx ranges from − 145 to 200 MPa and an initial compressive stress state with a value of about 60 MPa was observed.

Fig. 9
figure 9

Residual stress trend and measurements errors of AB 316L samples at z = 16 mm with different scanning strategies in terms of σx, σz and τxz

The comparison of σz of the samples obtained with different scanning strategies is illustrated in Fig. 9(b). C-0090-AB sample showed an initial tensile state of 100 MPa, and the stresses were characterized by a positive value for almost all the analyzed depth. At a depth of about 0.7 mm, a low compressive state was observed. After a depth of 0.8 mm, the value of the stress increased significantly and reached the maximum value of 240 MPa. C-0060-AB sample presented an initial compressive stress state with a tension value of about 50 MPa. Along the analyzed depth, the stress values range from − 80 to 160 MPa with the maximum tensile stress that was observed at a depth of about 0.8 mm.

These results highlight that C-0090-AB sample was characterized by a positive value of stress for almost all of the considered depth with a quite constant value. Instead, C-0060-AB sample showed an oscillatory sine-wave-like trend with a positive mean value. Comparing the graphs, it is possible to observe that with the 0°-90° deposition strategy (C-0090-AB sample) the residual stresses generated in the sample were lower than the ones developed using the 0°-60° deposition strategy (C-0060-AB sample) for almost all of the analyzed depth.

The shear stresses measured on the samples are illustrated in Fig. 9(c). It is possible to observe that the value of shear stress (τxz) was at least two times lower compared to the values of normal stresses (σx and σz). On C-0060-AB sample, the shear stress exhibited a slightly oscillatory trend between − 60 and 60 MPa; on the other hand, shear stress on C-0090-AB sample was characterized by an initial increasing trend, and after a depth of 0.1 mm, its value remained almost constant with a mean value of about 5 MPa.

All C-0060-AB residual stress profiles showed an oscillatory trend with a distance between peaks and valleys of about 0.2 mm. The curve of an as-built sample built with the 0°-60° strategy is reported in Fig. 10 together with the XY micrograph. The figure highlights a correspondence between microstructural features and stress values: In fact, it can be seen that the wavelength corresponds to the melt pool width. It is, however, difficult to define more specific correlations as the drill intersects several layers, and it is possible that the oscillation is due to a cumulative effect.

Fig. 10
figure 10

Micrographs of YZ cross section and residual stress trend of the C-0060-AB sample

C-0090-AB sample showed to have slightly lower stress values and to be characterized by more stable stress trends. Furthermore, the 0°-90° deposition strategy is more convenient from a production point of view as it facilitates the obtainment of geometrically stable parts. Because of these reasons, the effect of the distance from the substrate on the residual stress trends will be reported only for C-0090-sample. The lower and more stable residual stress values of 0°-90° samples can be attributed to the repeatable track length. The 0°-60° strategy is in fact characterized by tracks having very different lengths and therefore different cooling conditions (Ref 32).

Figure 11 depicts the residual stresses of the C-0090-AB sample measured at two different heights, which are located at z = 4 mm and z = 16 mm from the substrate, according to the scheme of Fig. 4. The components of σx recorded at different height are reported and compared in Fig. 11(a). In the upper position, near to the surface a tensile state occurred, in all the depth the stress ranges between − 30 and 230 MPa and a slightly oscillatory trend is visible. At the lower height of the C-0090-AB sample, near to the substrate, the stress exhibited a high positive value of about 640 MPa. Then, the stress abruptly decreases, and a compressive stress of 300 MPa was measured at a depth of about 0.1 mm. In both the analyzed heights, an oscillating trend of residual stresses was observed.

Fig. 11
figure 11

Residual stress trend and measurements errors measured on C-0090-AB sample at different heights in terms of σx, σz and τxz

The distribution of σz is depicted in Fig. 11(b). At the upper position, the stress showed an initial value of about 100 MPa; then, along the depth up to about 0.8 mm small fluctuations with an amplitude of about 50 MPa are observed. Then, for depth larger than 0.8 mm the stress sharply increased reaching the value of about 240 MPa. In the lower position, the stress showed an initial tensile state with a value of about two times higher compared to the one measured in the upper position. The maximum tensile stress of 570 MPa was measured at a depth of about 0.1 mm. After 0.1 mm, the stress value crossed the zero line and exhibited a compression state for almost all the analyzed depth.

From the results represented in Fig. 11, it was observed that for a depth lower than 0.1 mm the measured normal stresses were higher than the 80% of yield stress values of the tested material. As a consequence, according to ASTM E837-13a standard (Ref 26), in this case the evaluation of the residual stresses gives only a qualitative indication of the stress state.

In addition, analyzing the graphs in Fig. 11, it is possible to observe that the residual stresses measured in the lower position were higher than those measured in the upper position of the sample. This could be attributed to the lower cooling rate observed in the upper layers due to the increase in temperature induced by the process as demonstrated by previous works (Ref 1, 19).

Finally, Fig. 11(c) shows the comparison of shear stresses (τxz) measured on C-0090-AB sample. It was possible to observe that both in the upper position and in the lower position, the initial values of τxz were negative. In the upper position, the value of shear stress was almost constant; instead in the lower position, the shear stress was characterized by an oscillatory trend between − 15 and 70 MPa.

Heat Treatment

In order to reduce the residual stress and to homogenize the microstructure, two heat treatments were carried out on AB 316L samples. The XRD data of AB and HT samples built with the 0°-60° deposition strategy are reported in Fig. 12(a) and allowed to assess the effect of the heat treatments on the microstructure of LP-DED 316L samples. It is important to underline that equal results were obtained for the C-0090 samples, confirming that the deposition strategy did not affect the microstructure of the samples.

Fig. 12
figure 12

(a) XRD pattern of AB and HT samples and (b) inset of austenite (111) and ferrite (110) peaks

The XRD patterns of AB and HT samples reveal that in all cases the material is made of austenite (γ) and ferrite (Fig. 12). The ferrite peak intensity increases as a consequence of the heat treatments performed, as already observed by Saeidi et al. (Ref 30) on 316L samples produced by laser powder bed fusion (LPBF). Finally, it is also important to note that the gas atomized powder does not show the ferrite peak suggesting a different solidification condition of the processes.

Finally, Fig. 12(b) also demonstrates that there is a preferential orientation in the AB and HT samples indicated by the different intensity of the (111) austenite peak. This result suggests that the recrystallization of the microstructure did not arise at these temperatures.

The effect of the heat treatments on the hardness is reported in Fig. 13, revealing that 600 °C heat treatment does not influence the hardness values, while the 800 °C heat treatment causes a reduction in the hardness values. This reduction is probably due to the microstructure evolution and in particular to the reduced dislocation density and to the partial dendrite coarsening. Dendrite coarsening of AM 316L samples heat-treated at 600-800 °C was observed by several authors (Ref 33, 34).

Fig. 13
figure 13

Vickers hardness values of AB and HT samples

It is interesting to underline that the samples built with the 0°-90° deposition strategy always have a slightly higher hardness value, and the differences are, however, very low and generally lay within the standard deviation.

The homogenization effect of these heat treatments can also be observed by comparing the hardness trend along the building direction of AB and HT samples (Fig. 14). After 800 °C HT, the hardness values become almost constant along the building direction. This suggests that the post-processing heat treatments cause a modification of the microstructure and eliminate the effects, observed on AB samples, of the different cooling rates experienced at different heights. Unexpectedly, the C-0090-800 sample still shows a slight decreasing hardness trend along the Z direction, suggesting that in this case the homogenization was not completed.

Fig. 14
figure 14

Hardness trend along the Z direction of AB and HT samples built (a) with 0°-60° and (b) 0°-90° deposition strategy

The most representative tensile curves of AB and HT LP-DED 316L samples reported in Fig. 15 show that the heat treatment performed at 600 °C did not affect the tensile performances of the alloy. The YS value remains, in fact, constant, while there is a slight increase in the UTS values. A similar result was observed by Leuders et al. (Ref 35) on heat-treated L-PBF 316L samples.

Fig. 15
figure 15

Tensile curves of 316L LP-DED AB and HT samples

The heat treatment performed at 800 °C, on the contrary, affects the mechanical behavior of the samples; in fact, a strong reduction in the YS was observed. It is important to underline that these samples are also characterized by the highest elongation probably due to the initial coarsening of the dendritic structure (Ref 34). A similar phenomenon was observed by Morrow et al. (Ref 18) on LP-DED sample after 1 h at 1060 °C. Furthermore, the heat treatment performed at 800 °C causes an increase in the strain hardening phenomenon, probably because of the reduced dislocation density of these samples (Ref 35).

Table 2 summarizes the mechanical properties of the studied LP-DED 316L samples in the AB and HT conditions, compared with literature data of 316L AM samples. To the best of the author’s knowledge, only recrystallization heat treatments were performed on LP-DED 316L samples by Yadollahi et al. (Ref 6) and Morrow et al. (Ref 18). These studies observed a recrystallization of the microstructure and correlated the observed reduction in the YS and the UTS mainly with the microstructural coarsening and to the δ-ferrite loss. Furthermore, Yadollahi et al. (Ref 6) reported a lower variation in the YS due to the homogenization effect of the heat treatment.

Table 2 Tensile properties of LP-DED AISI 316L samples

A lower temperature heat treatment was performed by Leuders et al. (Ref 35) on LPBF 316L samples. The authors reported only a slight reduction in the YS probably due to the reduced dislocation density and confirmed that no recrystallization happened at this temperature for AM 316L samples.

The comparison of residual stress profiles measured on the lower position of samples obtained with the 0°-90° strategy and subjected to different heat treatments is reported in Fig. 15.

The comparison was reported only for the lower part of samples because, as shown in Fig. 11, this area is characterized by the presence of the highest residual stress and it is, therefore, more suitable to appreciate the residual stress reduction. Furthermore, this comparison is performed on samples built with the 0°-90° deposition strategy because, as previously discussed, it has more stable residual stress trends and it is more interesting from a processing point of view.

It is possible to observe that the values of the components of the stress profiles along the x-axis and z-axis on the HT samples are lower than ones measured for the AB sample.

In particular, Fig. 16 shows that the stresses value decreased more under high heat treatment temperature. In detail, the stresses observed in the C-0090-800 sample are almost nil for the analyzed depth. The treated samples showed a similar behavior along the x-axis and along the z-axis with an initial compressive state. After a depth of about 0.1 mm, the stress value slightly increased and a tensile state was observed for the almost all the investigated depth.

Fig. 16
figure 16

Residual stress trend and measurements errors of AB and HT samples at z = 4 mm in terms of σx, σz and τxz

Comparing the residual stress distributions depicted in Fig. 16, it is possible to note that the effect of heat treatments was more effective near the external surface in which a reduction in the stress value of about 300 MPa was recorded. However, slight oscillatory stress trends are observed for the analyzed samples and this means that the heat treatment does not affect the oscillatory nature of the stress distribution.

The comparison of shear stress (τxz) distribution on sample C-0090-AB, C-0090-600 and C-0090-800 is represented in Fig. 16(c). It was observed that lower shear stresses were obtained in the treated samples. In addition, it was possible to observe that the heat treatment temperature did not significantly influence the value of shear stress.

Conclusions

In the present work, the effect of the deposition strategy and the heat treatment on the microstructure, the properties and the residual stresses of LP-DED 316L samples was investigated. The results can be summarized as follows:

  • The deposition strategy does not strongly affect the microstructure and the hardness of LP-DED samples. In all samples, in fact, the microstructure was constituted of elongated grains containing fine austenitic dendrites surrounded by residual δ-ferrite. It was observed that the hardness values depend on the distance from the substrate showing higher values in the lower part due to the higher cooling rates experienced during the building process.

  • Residual stresses seem to be only slightly affected by the deposition strategy and the distance from the building platform.

  • The heat treatments performed caused the reduction and the homogenization of the hardness values which start to lose the Z dependence.

  • The residual stresses are strongly reduced by selected heat treatments. The heat treatments suggested in this work might be promising for LP-DED 316L SS. The 600 °C HT allows only a partial reduction in the residual stress, but keeps very high mechanical properties. The 800 °C HT, on the contrary, allows an almost complete reduction in residual stresses, but causes a 20% reduction in the YS, but a 8% increase in the elongation value. The first heat treatment might be therefore more suitable for the applications in which high strength is required, while the 800 °C one is more suitable when complex geometries are built and residual stress might have detrimental effects.