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npj Materials Degradation volume 8, Article number: 61 (2024 ) Cite this article Machining Service
In Type 316 L stainless steel fabricated from gas-atomized powder via spark plasma sintering, lack-of-fusion pores and MnS inclusions were identified as possible pitting initiation sites. Through potentiodynamic polarization with different working electrode areas, the distribution density of the identified pit initiation sites was compared with that of sites found on wrought Type 316 L stainless steel. Surface observations of the sintered Type 316 L after polarization suggest that pitting corrosion was initiated at a location where both MnS and pores existed. By reducing the porosity and removing MnS, the roles of pores and MnS inclusions in the initiation of pitting corrosion were investigated.
Powder metallurgy is a near-net-shape technology that provides metals with unique microstructures unobtainable by conventional melting and casting processes, making it both commercially useful and scientifically interesting1,2,3,4. Furthermore, metal additive manufacturing using powders has attracted attention for enabling rapid fabrication of complex shapes with high precision, and has been successfully applied to stainless steels5. However, sintered stainless steels fabricated using powder metallurgy techniques tend to have lower corrosion resistance than wrought stainless steels6,7. Otero et al.6 studied the corrosion resistance of Type 304 L and Type 316 L stainless steels fabricated via powder metallurgy in sulfuric and phosphoric acids, suggesting that the corrosion rates of sintered stainless steels were high because of the corrosion attack initiated in the pores by crevice corrosion. Pandya et al.7 investigated the effect of sintering temperature on the corrosion resistance of Type 316 L stainless steel in 0.05 M H2SO4. They demonstrated that the sintering temperature affects the porosity of sintered stainless steel—the higher the porosity, the lower the corrosion resistance. They concluded that the pore surface area added to the apparent surface area, thereby increasing the overall surface area exposed to corrosion. In chloride solutions at near-neutral pH, the pitting corrosion resistance of sintered stainless steel has also been reported to be lower than that of wrought stainless steel8,9. Even though this is attributed to the lack-of-fusion (LOF) pores, the detailed role of the pores in the pit initiation process remains unclear10,11,12,13,14.
In addition to pores, non-metallic inclusions15,16,17,18 are expected to play an important role in pit initiation. In particular, sulfide inclusions such as MnS are known to cause pitting corrosion in stainless steels19,20,21,22,23,24,25. Therefore, the influence of sulfide inclusions should be considered when discussing the pitting corrosion of sintered stainless steels. Laleh et al.26 investigated the localized corrosion of stainless steels produced by selective laser melting (SLM). They found the specimens containing large LOF pores with diameters larger than 50 µm were extremely susceptible to localized corrosion. However, no traces of MnS inclusions were detected due to rapid cooling by the SLM process. Other additive manufacturing techniques, such as binder jet printing, require relatively long sintering time and do not inhibit MnS formation. Therefore, it is essential to clarify the roles of pores and MnS inclusions in the initiation of pitting corrosion in order to improve the localized corrosion resistance of additive manufactured stainless steels.
In this study, we fabricated sintered Type 316 L stainless steel from gas-atomized powder. The distribution density of the pit initiation sites was investigated by varying the working electrode area from 5 × 5 mm to 300 × 300 µm. Surface observations of pit initiation sites were performed after potentiodynamic polarization. To clarify the roles of pores and sulfide inclusions in the initiation of pitting corrosion, the porosity and pore size were reduced by changing the sintering conditions, and the sulfide inclusions were removed by pulse electrolysis.
To fabricate sintered Type 316 L stainless steel (SS) from gas-atomized powder, spark plasma sintering (SPS) was performed at 1373 K for 20 min under a pressure of 30 MPa, and the resulting microstructure was compared to that of wrought Type 316 L SS. Figure 1 shows optical microscopy (OM) images of the specimen surfaces prepared by wet polishing using diamond pastes and alcohol-based lubricant. The left (Fig. 1a, c, and e) and right (Fig. 1b, d, and f) columns show OM images of wrought and sintered Type 316 L SS, respectively. Figures 1a and 1b show the as-polished surfaces. In wrought Type 316 L SS (Fig. 1a), almost no black dots were observed. In contrast, black dots were observed on the sintered Type 316 L SS (Fig. 1b). Figures 1c and 1d show the specimens after electrolytic etching in oxalic acid; enlarged images are shown in Figs. 1e and 1f. For wrought Type 316 L SS, grain boundaries were observed (Figs. 1c and 1e). For sintered Type 316 L SS, the grain boundaries were not etched, with many black dots appearing after etching. These black dots are thought to be the pores.
a, c, e Wrought Type 316 L SS; b, d, f sintered Type 316 L SS. a, b As-polished surface; c, d surface after electrolytic etching in 10 mass% oxalic acid; e, f enlarged view of (c, d), respectively. The surfaces were polished using a 1 µm diamond paste with an alcohol-based lubricant.
Figure 2a shows an OM image of a cross-section of sintered Type 316 L SS fabricated using an Ar-ion beam, with the black dots clearly visible. One of the black dots is indicated by the blue arrow. Figure 2b shows a backscattered electron (BSE) image of the area shown in Fig. 2a, and an enlarged BSE image of the area surrounded by the red lines in Figs. 2a and 2b is shown in Fig. 2c. Based on these images, the black dots were determined to be pores. A round shape resembling powder was observed at the position indicated by the red arrow; hence, the pore was associated with a lack of fusion. Comparing Figs. 1b and 2b, more pores were observed on the cross-section prepared by Ar ion beam polishing than on the surface prepared by wet polishing; pores were thought to be covered and hindered by polishing (see also a comparison of Figs. 1b and 1d). However, pores were present on the surface of the sintered Type 316 L SS. The number of pores was assessed from the observations of the cross-section of two different areas, and there were approximately 5 pores in a 100 × 100 µm area in the sintered Type 316 L SS fabricated in this study. To confirm the absence of segregation at the location of pores and around it, scanning electron microscopy (SEM) and energy-dispersive X-ray spectroscopy (EDS) analysis was conducted. Figure 2d shows the corresponding EDS maps in Fig. 2c. Although O was detected at the edge of the pores, no segregation of other elements was observed.
a OM image; b BSE images; c enlarged BSE image of the area surrounded by red lines in (a, b); d corresponding EDS maps of (c). The cross-section was fabricated with an Ar-ion beam. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
In this study, a commercially available wrought Type 316 L SS was used as the reference. Figure 3 shows an OM image, an SEM image, and the corresponding EDS maps of a typical inclusion in wrought Type 316 L SS. As shown in the image, the rolling direction of this specimen was vertical, which made the inclusion slightly elongated in the rolling direction. From the SEM images and EDS maps, the inclusion was divided into three parts: a round upper part, a middle part, and an oval lower part. O, Al, Si, and Ca signals were detected in the upper round part; S, Ca, and Cr signals in the middle part; and S, Ca, and Cr signals in the oval lower part. EDS point analysis was performed at Points 1–3, and the results are listed in Table 1. Si and Ca were detected in the area where O was present (upper section). In addition, Al was detected at the center of the region where O was present. In contrast, Ca was detected in the S-containing areas (middle and lower parts). Based on the above results, the oxide inclusions were thought to be complex oxide inclusions of Si and Ca with Al oxide nuclei (shown in the upper part of the SEM image), and the sulfide was determined to be CaS (shown in the middle and lower parts of the SEM image). The optical microscopy revealed that the number of inclusions larger than 1 µm was approximately 1 per 100 × 100 µm area. Based on the EDS analysis, the probability of S-containing inclusions was determined to be 80%, indicating that eight out of ten were sulfides. These are the average values observed and measured in four different areas. In this study, a commercial wrought Type 316 L SS with low S concentration (0.0009 mass%) was used, and CaS was formed, suggesting that Ca-treatment was applied. Ca treatment is a method of adding Ca to stainless steel; it can modify MnS to CaS and lower the S content of the steel matrix27,28,29,30. Ca- and Al-containing inclusions have been reported to react with S in the steel matrix during heat treatment to form Ca- and Al-oxides + CaS30, and it is presumed that a similar reaction occurred in the wrought Type 316 L assumed to be the Ca-treated steel used in this study. In contrast, the formation of MnS with oxide inclusions has been reported for austenitic stainless steels without Ca-treatment31.
a OM image; b SEM image; c–j corresponding EDS maps.
Figure 4 shows an OM image, SEM image, and corresponding EDS maps of a typical inclusion in sintered Type 316 L SS. The inclusions were divided into four areas based on the SEM images and EDS maps, as shown in Fig. 4a. O and Si signals were detected in Area 1. S, Mn, and Cr were detected in Area 2. O, Si, Cr, and Fe were detected in Area 3. O and Si were detected in Area 4. EDS point analysis was conducted at Points 4–6, and the results are listed in Table 1. Si, Cr, and Fe were also detected in areas where O was detected (Areas 1, 3, and 4). In contrast, Mn was detected in the S-containing area (Area 2). Based on the above results, the oxide inclusions were determined to be complex oxides of Si, Cr, and Fe, whereas the sulfide was determined to be MnS. The optical microscopy revealed that the number of inclusions larger than 1 µm was approximately 6 per 100 × 100 µm area. Based on EDS analysis, the probability of S-containing inclusions was determined to be 30%, indicating that three out of ten were sulfide inclusions. These are the average values for the four different areas.
a OM image; b SEM image; c–i corresponding EDS maps. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
To assess the pitting corrosion resistance of the specimens and to ascertain the distribution density of pit initiation sites, potentiodynamic polarization measurements were performed with an electrode area of 5 × 5 mm to 300 × 300 µm. Figure 5a shows the potentiodynamic polarization behavior of wrought Type 316 L SS. In all cases, polarization measurements were started from −0.3 V, and anodic current was observed above approximately −0.1 V. When the electrode area was 5 × 5 mm, the polarization curve showed passive behavior in the potential range from 0 to 0.6 V; however, current density sharply increased at 0.6 V because of pit initiation. In the cases of 1 × 1 mm and 500 × 500 µm areas, although a gradual increase in current density was observed above 1.1 V due to the oxygen generation reaction, no rapid increase in current density occurred until 1.4 V. Figure 5b–d show surfaces of the wrought Type 316 L SS after polarization. On the 5 × 5 mm specimen, a large pit in the center of the electrode area and some small pits were observed. In the cases of 1 × 1 mm and 500 × 500 µm electrodes, although crevice corrosion was observed at the boundary between the electrode area and masking, no pitting corrosion was found to occur in the passive region. The distribution density of pit initiation sites on wrought Type 316 L SS was expected to be one or less per 1 × 1 mm.
a Polarization curves for the electrode area of 5 × 5 mm, 1 × 1 mm, and 500 × 500 μm; b–d electrode surfaces after polarization with electrode areas of 5 × 5 mm, 1 × 1 mm, and 500 × 500 μm, respectively.
The wrought Type 316 L SS used in this study contained CaS, and the pit initiation sites were thought to be CaS32,33,34. In the polarization measurements of wrought Type 316 L SS, the pitting potential increased with decreasing electrode area, a tendency consistent with that reported by Burstein et al.35,36. In addition, the dissolution of sulfide inclusions causes pitting corrosion in stainless steels21,37,38. Chiba et al. polarized a small area with and without sulfide inclusions in NaCl solutions. They reported that pitting corrosion did not occur in areas without sulfide inclusions39. It was revealed that stainless steel matrices intrinsically exhibit high pitting corrosion resistance. Moreover, pitting corrosion is known to occur readily in large sulfide inclusions22,40. Assuming that sulfide inclusions were randomly distributed on the specimen surfaces and that pitting corrosion was initiated at large sulfide inclusions, the pitting potential increased with decreasing working electrode area. Furthermore, when the electrode area becomes sufficiently small, no initiation of stable pitting occurs, even in chloride solutions41.
Figure 6a shows the potentiodynamic polarization behavior of sintered Type 316 L SS. Polarization was also started from −0.3 V. With the electrode areas of 5 × 5 mm and 1 × 1 mm, the anodic current density significantly increased above the corrosion potential. In the cases of 500 × 500 µm and 300 × 300 µm electrodes, although current spikes were generated, the passive region was certainly observed. The 500 × 500 µm and 300 × 300 µm electrodes showed a large and rapid increase in current density at 0.5 V and 1.0 V, respectively. Figure 6b–e shows the surfaces of the specimens after polarization. Several pits were observed in the 5 × 5 mm and 1 × 1 mm electrodes. Therefore, the large increase in current density above the corrosion potential was not caused by active dissolution but rather by pitting corrosion induced by chloride ions42,43. For the 1 × 1 mm electrode, the number of pits was lower than that of the 5 × 5 mm electrode; however, some large pits were generated. For the 500 × 500 µm and 300 × 300 µm electrodes, crevice corrosion was observed at the boundary between the masking and electrode areas, but no large pits were seen. For the sintered specimens, there was at least one site per 1 × 1 mm area where pitting corrosion occurred. However, more initiation sites were expected because micropitting was reported to be the initial stage of crevice corrosion44,45,46 and crevice corrosion occurred in the passive region for both 500 × 500 µm and 300 × 300 µm. In this study, the distribution density of pit initiation sites on sintered Type 316 L SS was found to be much higher than that on wrought Type 316 L SS.
a Polarization curves for the electrode area of 5 × 5 mm, 1 × 1 mm, 500 × 500 μm, and 300 × 300 μm; b–e electrode surfaces after polarization with electrode areas of 5 × 5 mm, 1 × 1 mm, 500 × 500 μm, and 300 × 300 μm, respectively. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
Pitting corrosion of stainless steels is known to be caused by the breakdown of passive film by chloride ions. Because the interaction between chloride ions and passive film occurs at atomic scale42,43, the pitting corrosion resistance is expected to be independent of electrode area. However, for the sintered Type 316 L SS, the pitting potential increased with decreasing working electrode area, like the wrought Type 316 L SS. This suggests that the stainless steels used in this study have very low corrosion resistance regions activated by chloride ions. As mentioned above, sulfide inclusions existed in both sintered and wrought Type 316 L SS: sintered Type 316 L SS contained MnS and wrought Type 316 L SS contained CaS. As discussed below, the differences in the pitting corrosion resistance and distribution density of the pit initiation sites cannot be attributed to the differences in the type of inclusions. In the sintered Type 316 L SS, the number of inclusions larger than 1 µm was approximately 6 per 100 × 100 µm area, as mentioned earlier. The probability of S-containing inclusions was 30%. Therefore, approximately 180 MnS inclusions with dimensions of 1 × 1 mm were obtained. On the other hand, in the wrought Type 316 L SS, optical microscopy revealed that the number of inclusions larger than 1 µm was approximately 1 per 100 × 100 µm area, and among them, the probability of S-containing inclusions was determined to be 80% based on EDS analysis. Thus, approximately 80 CaS inclusions with dimensions of 1 × 1 mm were obtained. Thus, the distribution density of sulfide inclusions on the wrought and sintered specimens was 80:180. However, not all inclusions undergo pitting. Moreover, it has been reported that CaS corrodes more readily than MnS32,33. Therefore, assuming that the pit initiation sites on the sintered Type 316 L SS were MnS alone, the pitting corrosion resistance was expected to be roughly equivalent to or slightly inferior to that of the wrought Type 316 L SS. However, the pitting corrosion resistance of sintered Type 316 L SS was significantly lower than that of wrought Type 316 L SS. For example, in the case of the 1 × 1 mm electrode area, a large current was observed for sintered Type 316 L SS in the passive region of wrought Type 316 L SS, and many pits were generated (see Fig. 6). Other than the inclusions, the lack-of-fusion pores are likely to be a possible factor that reduces pitting corrosion resistance.
To ascertain the initiation site of pitting on the sintered Type 316 L SS, detailed surface observations were performed on an electrode area of 1 × 1 mm. Figures 7a and 7b show OM images of a pit initiation site (surrounded by the red lines in Fig. 6c) before and after polarization, respectively. In Fig. 7a showing images before polarization, a black dot is indicated by the red arrow. The black dots represent inclusions or pores. After polarization, as shown in Fig. 7b, brown corrosion products can be seen around the area where the black dot exists, and dark streak patterns were observed, which seemed to be dissolved areas (corroded areas). No round pits were observed on the sintered Type 316 L SS. This suggests that the initiation and growth of pitting corrosion did not simply occur due to the breaking down of the passive film or dissolution of inclusions.
a Before and b after polarization.
Figure 8a–j shows an OM image, a BSE image, and the corresponding electron probe microanalyzer (EPMA) maps of the area shown in Fig. 7b. As shown in Figs. 8a and 8b, the corroded areas that appear dark in the OM image do not correspond to the black areas in the BSE image. This was likely due to the fact that the BSE image shows holes existing below the surface because of electrons migrating from the surface to the interior of the specimen. From the EPMA maps, O was entirely detected in the corroded areas. However, Si was detected at specific positions that are thought to be oxide inclusions. Additionally, S was detected at the same positions as Mn in the corroded areas, indicating the presence of MnS inclusions. The results of the EDS analysis at Points 7–9 are listed in Table 2. The relative concentrations of Mn and S at Points 7 and 8 were higher than those in the steel matrix (Point 9).
a OM image; b BSE image; c–j corresponding EPMA maps. k, l SEM image and the corresponding EDS maps of cross-sections at the sites marked by the red lines in (b). The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
Figures 8k and 8l show SEM images and corresponding EDS maps of the cross-sections at the sites marked by the red lines in Fig. 8b. The cross-sections were fabricated using focused ion beam milling. As shown in Fig. 8k, holes were observed below the surface. Part of the wall on the right side of the image appeared relatively smooth, suggesting that it was a lack-of-fusion pore. Moreover, as shown in Fig. 8l, a large hole was observed under the surface, confirming that pitting corrosion occurred. Because S and Mn were detected at the opening of the pit with a large internal cavity, both MnS and lack-of-fusion pores were likely involved in the initiation of pitting corrosion in sintered Type 316 L SS. In other words, pitting corrosion of sintered Type 316 L SS is thought to occur when MnS and pores exist at almost the same location.
To further clarify the role of pores in pit initiation, the SPS conditions were changed to reduce the number and size of pores. Figure 9 shows OM images of the sintered Type 316 L SS surfaces. The left (Fig. 9a, c, e, g) and right (Fig. 9b, d, f, h) columns, respectively, show OM images of the as-polished specimens and the specimens after electrolyte etching. The specimens shown in Figs. 9a and 9b (Specimen A) were sintered at 1373 K for 20 min at 30 MPa and used in the experiments described above. By contrast, Figs. 9c and 9d (Specimen B) show the case in which the holding time was extended to 40 min. Figures 9e and 9f (Specimen C) show the results when the annealing temperature was increased to 1423 K, and Figs. 9g and 9h (Specimen D) showed the results when the pressure was increased to 40 MPa. As shown in Fig. 9, pores were clearly observed after electrolytic etching. From the images of the etched surfaces, the number of pores, area ratio of the pores, and average area of the individual pores were measured using image analysis software. The results are summarized in Table 3. To reduce porosity and pore size, the most effective SPS parameter was determined to be the sintering pressure.
a, b Specimen A (1373 K, 20 min, 30 MPa); c, d Specimen B (1373 K, 40 min, 30 MPa); e, f Specimen C (1423 K, 20 min, 30 MPa); and g, h Specimen D (1373 K, 20 min, 40 MPa). a, c, e, g As-polished surfaces; b, d, f, h etched surfaces.
Figure 10 shows the polarization behavior of Specimens A–D. The polarization curve of wrought Type 316 L SS is presented as a reference. Specimens B–D, in which the porosity and pore size were reduced, exhibited lower current densities than Specimen A. This indicates that the reduction in porosity and pore size was effective in improving the pitting corrosion resistance. From this result, there is no doubt that pores are involved in the initiation of pitting corrosion in sintered Type 316 L SS.
Polarization curves for working electrode area of 5 × 5 mm.
To assess the effect of sulfide inclusions on pitting corrosion in sintered Type 316 L SS, the MnS inclusions were removed. The removal (dissolution) of sulfide inclusions is an effective way to improve the pitting corrosion resistance of stainless steels21. Immersion in nitric acid is a commonly used approach that is effective in removing sulfides and forming Cr-rich surface oxide films47,48,49,50. Potentiodynamic51, potentiostatic52, and cyclic polarization treatments53,54 have been reported to be effective. Acidic solutions are generally used in immersion and polarization processes; however, they may dissolve the steel surface along with inclusions, and the effects of passive film modification on pitting corrosion resistance cannot be eliminated. From this perspective, solutions with near-neutral pH are desirable. However, as sulfides do not dissolve easily at a near-neutral pH, a high voltage should be applied to dissolve the steel matrix. Therefore, in this work, we focused on high-voltage pulse electrolysis (PE). As PE-treatment, a voltage of ±20 V was applied in 0.1 M NaNO3 (pH 6.0) at 293 K.
Figure 11 shows the polarization curves of the PE-treated sintered Type 316 L SS in 0.1 M NaCl. SPS was performed at 1373 K for 20 min at a pressure of 30 MPa. These sintering conditions were those under which a large number of pores existed (Fig. 9a, b). The size of working electrode was 5 × 5 mm. For reference, the polarization curves of untreated sintered Type 316 L SS and untreated wrought Type 316 L SS are provided. The PE-treated Type 316 L SS specimen exhibited higher pitting corrosion resistance than the untreated specimens. In the case of the untreated sintered Type 316 L SS, a large current was observed even at lower potentials owing to pitting (see Fig. 6a, b); however, the current density of the PE-treated specimen drastically decreased. Furthermore, the pitting potential of the PE-treated sintered Type 316 L SS was higher than that of the untreated wrought Type 316 L SS, although many current spikes were observed in the passive region.
Polarization curves for working electrode area of 5 × 5 mm. Polarization curves of untreated sintered and untreated wrought Type 316 L SSs are presented as a reference. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
In stainless steels, the surface oxide films affect the pitting corrosion resistance42,43,55,56,57; therefore, the effect of PE-treatment on the oxide films was analyzed. Figure 12 shows the auger electron spectroscopy (AES) depth profiles of the PE-treated and untreated specimens. The thicknesses of the oxide films on the PE-treated sintered Type 316 L SS, untreated sintered Type 316 L SS, and untreated wrought Type 316 L SS were 2.4, 1.9, and 1.9 nm, respectively. The oxide film on the PE-treated sintered Type 316 L SS was slightly thicker than that on the other specimens; however, this factor was unlikely to significantly change the pitting corrosion resistance. No compositional differences were observed between the treated and untreated specimens. In particular, no significant difference was observed with respect to Cr, which is an important factor in corrosion resistance. Therefore, the improvement in the pitting corrosion resistance by the PE-treatment could not be attributed to the modification of the passive film.
AES depth profiles of a PE-treated sintered Type 316 L SS; b untreated sintered Type 316 L SS; and c untreated wrought Type 316 L SS. The PE-treatment was performed at 293 K. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
To confirm the effect of the PE-treatment on inclusion dissolution, SEM/EDS analysis was performed. Figure 13 shows an SEM image and the corresponding EDS maps of an inclusion on the PE-treated sintered Type 316 L SS. As shown in the SEM image, this inclusion was divided into two parts: upper and lower. In the lower part, the boundary between the steel matrix and the inclusion was dissolved. In the EDS maps, O and Si signals were detected in the upper part, whereas O, S, and Mn signals were detected in the lower part. EDS analysis was conducted at Points 10–12 and the results are listed in Table 4. It was confirmed that the upper part was composed of Si oxide and the lower part of MnS. Therefore, it was concluded that the dissolved part was the MnS/steel boundary. In addition, an O signal detected in the lower part confirmed the formation of an oxide layer on the MnS inclusion. As a result of the pulse electrolysis of the sintered Type 316 L SS, the MnS/steel boundary was dissolved. According to previous reports39,58,59,60,61,62,63, pitting corrosion occurs at the MnS/steel boundary; therefore, the removal of these pit initiation sites by pulse electrolysis is thought to be effective in suppressing the initiation of pitting in sintered Type 316 L SS. In this way, it was found that the corrosion resistance was greatly improved even without entirely dissolving MnS, and that MnS played a major role in the initiation of pitting corrosion of sintered Type 316 L SS.
a SEM image; b–g the corresponding EDS maps. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
To further clarify the role of MnS in pit initiation, PE-treatment was performed at a higher temperature of 353 K to completely dissolve MnS. SPS was also performed at 1373 K for 20 min under a pressure of 30 MPa, and the resulting specimen contained numerous pores (Fig. 9a, b). Figure 14 shows the SEM image and corresponding EDS maps of a typical S-containing inclusion before and after the PE-treatment. To eliminate the effect of electron irradiation on the PE-treatment, the specimen surface was gently polished using 1 µm diamond paste for 20 s before the PE-treatment. As seen in this figure, the signals of S and Mn disappeared after PE-treatment at 353 K. Figure 15 shows the polarization behavior in 0.1 M NaCl. For reference, the results for the PE-treated specimens at 293 K are presented. In the case of sintered Type 316 L SS PE-treated at 353 K, the number of current spikes due to metastable pitting was drastically reduced, and no stable pitting was generated. Even after PE-treatment at 353 K, pores were certainly present, but the pitting corrosion resistance was greatly improved by the removal of MnS. This indicates that MnS plays a critical role in pit initiation in sintered Type 316 L SS.
SEM images and the corresponding EDS maps: a before and b after PE-treatment. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
Polarization curves for working electrode area of 5 × 5 mm. Polarization curves of sintered Type 316 L SS PE-treated at 293 K is presented as a reference. The SPS was performed at 1373 K for 20 min under a pressure of 30 MPa.
The mechanism of pit initiation on sintered Type 316 L SS is as follows. During anodic polarization in 0.1 M NaCl, the pores whose openings were covered with steel by polishing were exposed to the solution owing to the dissolution of MnS. Because the pores are deep and narrow, it is considered that corrosive species (Cl−, H+, and S-species) tend to accumulate in the pores and become stable pits. Small sulfide inclusions that would have been metastable pits in the wrought Type 316 L SS were presumed to be stable pits in the sintered Type 316 L SS owing to the co-existence of pores and MnS. Although the pores could not be removed by pulse electrolysis, MnS could be removed. Thus, pulse electrolysis is thought to improve the pitting corrosion resistance. In other words, if only pores without MnS were present, the possibility of pitting would be extremely low. In sintered stainless steels, reducing sulfide inclusions is effective for improving pitting corrosion resistance. MnS inclusion is generally inhomogeneous both in structure and chemical composition. Zheng et al.31 found that nano-sized MnCr2O4 crystals were embedded in MnS inclusions. They demonstrated that MnS initially dissolves at the MnCr2O4/MnS interface because of generating local MnCr2O4/MnS nano-galvanic cells. Therefore, the boundary between oxide and sulfide in the inclusions may be the initiation site of inclusion dissolution in sintered Type 316 L SS. However, the details are unclear, and further research is needed on the galvanic coupling between the oxide and sulfide.
In this study, the initiation sites of pitting on sintered Type 316 L SS were investigated using wrought Type 316 L SS as a reference. Moreover, by reducing the porosity and removing MnS, the roles of pores and MnS inclusions in the initiation of pitting corrosion were investigated. The following conclusions were drawn.
1. There were many pores in the sintered Type 316 L SS, which were determined to be lack-of-fusion pores in terms of morphology. No segregation was observed around the pores. MnS was observed in the sintered Type 316 L SS, while the wrought Type 316 L SS contained CaS inclusions.
2. The pitting corrosion resistance of sintered Type 316 L SS was significantly lower than that of wrought Type 316 L SS. In potentiodynamic polarization in 0.1 M NaCl, when the working electrode area was 1 × 1 mm or larger, many pits were generated, and no passive region was observed on the polarization curves of the sintered Type 316 L SS.
3. EPMA analysis revealed the existence of MnS at the pit initiation sites on sintered Type 316 L SS. Moreover, based on the pit morphology, both MnS inclusions and lack-of-fusion pores were likely involved in the initiation of pitting corrosion.
4. By reducing the porosity and pore size, pitting corrosion resistance of the sintered Type 316 L SS was slightly improved.
5. By pulse electrolysis in 0.1 M NaNO3, the boundaries of MnS/steel dissolved at 293 K, and MnS dissolved entirely at 353 K. In both cases, the pitting corrosion resistance significantly improved.
Sintered Type 316 L SS was fabricated by SPS. Gas-atomized Type 316 L SS powders were used as the raw materials, and their chemical compositions are listed in Table 5. Prior to SPS, the Type 316 L SS powder (7 g) was packed into a graphite die with a diameter of 15 mm. To prevent air contamination (oxygen and nitrogen), the packing was performed in a glove box with Ar gas. In the SPS process, the graphite die with the powder was heated from room temperature to the sintering temperature (1373 or 1423 K) for 20 min and then held for 20 or 40 min under vacuum. The current was approximately 820 A and the pressure was set to 30 or 40 MPa. Subsequently, the current was cut off and the specimen was cooled to room temperature in an SPS vacuum chamber. The SPS specimens were cylindrical with a diameter of 15 mm and a thickness of 5 mm. After sintering, the carbon contamination on the cylindrical specimen surfaces was removed, and solution treatment was conducted at 1373 K for 30 min (water-quenched). Subsequently, the cylindrical specimens were sliced to produce two disc-shaped specimens of approximately 2.5 mm thickness, and the sliced surfaces were polished down to 1 µm using diamond pastes. An alcohol-based lubricant (DP-Lubricant Blue) was used to polish the diamond paste.
Commercially available wrought Type 316 L SS with a thickness of 2 mm was used as the reference material. The chemical composition of wrought Type 316 L SS is listed in Table 5. Wrought Type 316 L SS was cut into 15 × 25 mm coupons and solution-treated at 1373 K for 30 min (water quenched). The specimen surface was polished down to 1 µm using diamond pastes and lubricant (DP-Lubricant Blue).
Polarization measurements were conducted in 0.1 M NaCl at pH 6.0. The pH of the solution was adjusted using NaOH. All the chemicals used in this study were of analytical grade, and deionized water was used to prepare the solutions.
Potentiodynamic polarization measurements were performed in naturally aerated 0.1 M NaCl at 298 K. The scan rate was set at 23 mV min−1 and the conventional three-electrode method was used. The reference electrode was Ag/AgCl (3.33 M KCl), and a Pt sheet served as the counter electrode. The working electrode area was varied from 5 × 5 mm to 300 × 300 µm to investigate the distribution density of pit initiation sites. Except for the working electrode area, all surfaces of the specimens were covered with resin. The area was measured accurately after polarization and a battery-powered potentiostat63 was used to reduce the electrical noise when the electrode area was less than 1 × 1 mm. All electrode potentials in this study are referred to as the Ag/AgCl (3.33 M KCl) electrode (0.206 V vs. SHE at 298 K).
PE was performed in 0.1 M NaNO3 (pH 6.0) using a two-electrode setup. A Pt sheet was used as the counter electrode. The working electrode was disk-shaped with a diameter of 8 mm. The pulse voltage was set to ±20 V, and the pulse frequency was 1 kHz. The pulse width was set to approximately 0.9 µs. Electrolysis was conducted for 10 min at 293 or 353 K. The specimens were then rinsed with deionized water and dried under a stream of N2. Subsequently, the specimens were stored in air for approximately 12 h, and their pitting corrosion resistances were assessed using potentiodynamic polarization.
To observe the microstructure of specimens, electrolytic etching was conducted in 10 mass% oxalic acid at 2 A cm−2 for 20 s. Further, a cross-section preparation device (JEOL Cross Section Polisher) was used to analyze the cross-section of sintered specimens. The specimens were milled for 4 h at 6 kV using an Ar-ion beam. The as-polished and etched surfaces of the specimens were observed under an optical microscope (OM). A field-emission scanning electron microscope (FE-SEM) equipped with an EDS system was used to observe the surface and/or cross-section of the specimens. The acceleration voltage was set to 20 kV. It was difficult to distinguish between the S and Mo signals using EDS analysis. Therefore, an EPMA was used to separate the Mo and S signals. The electron acceleration voltage of EPMA was set to 15 kV. Additionally, focused ion beam (FIB) milling was used to create a cross-section of the specimen. Cross-sections were prepared by milling the specimen surface with a gallium ion beam and observed at a tilt angle of 30° using SEM.
AES and depth profiling were performed to analyze the surface oxide films on the specimens. The electron accelerating voltage was 3 kV and the beam current was 5 nA. The analyzed area was ca. 8 × 8 μm. The surfaces of the specimens were etched using an Ar+ beam with an accelerating voltage of 2 kV and a tilt angle of 30°. The etching rate was estimated to be 1.13 nm min−1 based on a thermally oxidized SiO2 film on Si. The full width at half maximum of the oxygen profile was used as the thickness of the surface oxide film.
The datasets used and/or analyzed in the current study are available from the corresponding author upon reasonable request.
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This study was supported by JSPS KAKENHI (grant numbers JP22H00254 and JP22KJ0240). The authors gratefully acknowledge support from the Ihara Science Nakano Memorial Foundation. The first author (H. Saito) is supported by a Grant-in-Aid for JSPS Research Fellows (Grant No. JP22J11168).
Department of Materials Science, Graduate School of Engineering, Tohoku University, 6-6-02, Aza-Aoba, Aramaki, Aoba-ku, Sendai, 980-8579, Japan
Haruka Saito, Masashi Nishimoto & Izumi Muto
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H.S.: investigation, methodology, data curation, visualization, writing—original draft, funding acquisition; M.N.: conceptualization, methodology, writing—review and editing; I.M.: conceptualization, methodology, writing—review and editing, supervision, and funding acquisition.
Correspondence to Haruka Saito or Izumi Muto.
The authors declare no competing interests.
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Saito, H., Nishimoto, M. & Muto, I. Pitting corrosion characteristics of sintered Type 316 L stainless steel: relationship between pores and MnS. npj Mater Degrad 8, 61 (2024). https://doi.org/10.1038/s41529-024-00482-6
DOI: https://doi.org/10.1038/s41529-024-00482-6
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