## 3 Results and Discussion
This density value is consistent with the studies in the literature [30]. Additionally, some twinning was observed in 316 LLS. These twins are annealing twins and occur during the recrystallization of close-packed structures [31]. Another factor influencing porosity formation is the physical deformation of the parts during the binder removal process [32]. Figure 7e shows the deformations in the samples obtained after high-speed and low-speed binder removal processes. It is clearly evident that an increase in the rate of temperature rise during the removal of the binder will lead to an increase in the size of the blister. The SEM images of the samples are provided in Figure 7c,d. In Figure 7c, the pores at the grain boundaries are more pronounced. Although both samples underwent binder removal and sintering processes at the same temperature, the binder removal process in the samples shown in Figure 7a,c was completed 25 times faster due to the higher speed at which the binder removal temperature was reached. Even though the binder removal temperatures were low, it is known that as time increases, diffusion also increases [33]. Therefore, the second sample (Figure 7b-d) had more time for grain boundary diffusion to occur. This is one of the primary reasons for the difference in porosity between the two samples. Furthermore, more intergranular porosity is observed in the second sample. In the first sample, the pores are more likely to be the inter-particle pores in the green parts, and it was concluded that the diffusion during sintering was not sufficient to close these pores. Due to the limited solubility of sulfur in austenite, it generally precipitates as sulfides (MnS or FeS) at the grain boundaries [34]. Particularly, the component observed in Point 4 of Figure 7d is believed to be a sulfur compound (Table 1). Despite the high oxygen composition at the grain boundaries in both samples, the oxygen content is higher in the first sample compared to the second (Table 1). This can be attributed to the difference in expansion between metal and oxide during the binder removal process, which took place over a longer duration in the second sample (316 LLS). This caused the brittle oxide layer on the particle surface to crack, and neck formation for diffusion began at the broken  
Figure 6: TGA analysis results of the 316L stainless steel sample. points [35]. Additionally, the EDS analysis taken from the grain interior regions confirms that the grains are in the austenite phase. Figure 8 shows the XRD graphs of green parts and sintered parts. It is clearly observed that the phases in the structure consist of a face-centered cubic (FCC) austenite phase and a body-centered cubic (BCC) ferrite phase (8), with no other phases detected. However, the XRD peak widths are broader in the sintered samples compared to the green parts. This is due to residual stresses formed during the heat treatment process in the parts [36]. Furthermore, the austenite phase is more prevalent in the 316 LLS sample compared to the 316 LHS sample. This can be attributed to the longer duration for which the sample was held at the austenitizing temperature. Figure 7: Microstructure images after the debinding process. (**a**) Optical microstructure images of 316 LHS; (**b**) optical microstructure images of 316 LLS; (**c**) SEM images of 316 LHS; (**d**) SEM images of 316 LLS and EDS analysis points; (**e**) effect of binder removal rate on blister height. The hardness values of the samples are shown in Figure 9. The sample labeled 316 LLS exhibited a higher hardness value (199 \(\pm\) 12) compared to the 316 LHS sample (149 \(\pm\) 5). A similar trend was observed in the density values of the parts. Additionally, the higher amount of austenite phase detected in the 316 LLS sample is another reason for the higher hardness of this part. As is known, the hardness of the FCC-structured austenite phase is higher than that of the BCC-structured ferrite phase [37].