Thickening of S-Phase and Sα-Phase of Various Stainless Steels Treated by Low Temperature Plasma Nitriding Using Screen

Abstract

This study aimed to observe the thickening of the S-phase and Sα-phase of various stainless steels subjected to low-temperature direct current plasma nitriding using screen (S-DCPN). Austenitic stainless steel SUS304, ferritic stainless steel SUS430, and duplex stainless steel SUS329J4L were treated using two different screens for comparison, namely, a Ni screen and a steel plate cold commercial (SPCC) screen. Plasma nitriding was performed at 673 K for 300 min under a 75% N₂ + 25% H₂ atmosphere at 100 Pa pressure of the mixed gas. After nitriding treatment, the samples were examined using X-ray diffraction (XRD) and glow discharge optical emission spectrometry (GD-OES), and their cross-sectional microstructure and surface microstructure were examined using an electron probe micro analyzer (EPMA). Nitrided samples were also subjected to Vickers hardness and pitting corrosion tests. Examination of the SUS304 samples revealed thickening of its S-phase and higher surface hardness and pitting corrosion resistance when nitriding was done with the Ni screen. This was due to excess nitrogen diffusion into the sample due to presence of the Ni screen than with the SPCC screen. In the SUS430 samples, thickening of the Sα-phase was not be observed. When the Ni screen was used during nitriding, higher surface hardness and less pitting corrosion resistance of sample were observed, along with enhanced nitrogen diffusion than the SPCC screen was used. In the SUS329J4L samples, upon nitriding with Ni screen, thickening of the S-phase was observed and surface hardness and pitting corrosion resistance of the sample were higher, which was attributed to the enhanced nitrogen diffusion into the sample than when nitriding with the SPCC screen.

 

This Paper was Originally Published in Japanese in J. Japan Inst. Met. Mater. 86 (2022) 62–70.

Fig. 2 GD-OES nitrogen profiles of (a) SUS304, (b) SUS430 and (c) SUS329J4L samples treated by S-DCPN using Ni screen and SPCC screen.

1. Introduction

In recent years, low-temperature nitriding has emerged as an important approach for the surface modification of stainless steels. It has been reported that nitriding of austenitic stainless steels above 723 K improves their surface hardness and wear resistance through the precipitation of chromium nitrides, but decreases the corrosion resistance due to a decrease in the amount of Cr solid solution in the matrix.^1–7) However, low-temperature nitriding performed at temperatures below 723 K has been reported to improve the surface hardness, wear resistance, fatigue resistance, and corrosion resistance by forming an expanded austenite (γ_N or S) phase with supersaturated nitrogen in the austenite (γ-Fe) phase of the base material.^1–13) Below 723 K, the diffusivity of Cr solid solution in the matrix is low and the nucleation of chromium nitride is suppressed, which mitigates the decrease in corrosion resistance due to chromium nitride precipitation.^4,12) Compared with the number of reports on low-temperature nitriding for austenitic stainless steels, there are relatively few reports on ferritic and duplex stainless steels. Low-temperature nitriding of ferritic stainless steels can form an expanded ferrite (α_N or Sα) phase, which is composed of a ferrite (α-Fe) phase with supersaturated N. For duplex stainless steels, the S- and Sα-phases are formed because the γ- and α-phases are mixed at an approximately 1:1 ratio; alternatively, the α-phase is transformed to the S-phase by the diffusion of nitrogen (γ stabilizing element) and only the S-phase is formed.^4,6,13–22) The mechanical properties and corrosion resistance of stainless steels can be improved by the formation of these expanded phases; however, a long processing time is required to achieve these properties because of the low-temperature nature of the treatment.

In general, there are three nitriding methods: gas nitriding using ammonia gas, salt-bath nitriding using cyanides, and plasma nitriding using plasma formed by glow discharge under a low vacuum containing nitrogen gas. Among these methods, plasma nitriding is more suitable than other methods for nitriding stainless steels because it can reduce and remove oxides, such as passive films on a material surface, through the collision of hydrogen ions when mixing hydrogen gas. In direct-current plasma nitriding (DCPN), which is a conventional plasma nitriding method, plasma is formed on a material surface by applying a cathodic voltage to the material. However, due to the high applied cathodic voltage, melting of the material by abnormal discharge and non-uniform nitriding treatment owing to the edge effect may occur.^23–28) Active screen plasma nitridation (ASPN), also known as cathodic cage plasma nitriding (CCPN), was developed to address these issues.^{23–25,27–40)} In ASPN, a material is insulated and voltage is applied only to a metal screen that is placed around the material in order to form plasma on the screen surface. Using this method, edge effects and material melting due to abnormal discharges are avoided. Various theoretical mechanisms have been proposed for ASPN. According to the “sputtering and deposition” theory proposed by some researchers,^25,30,34,35) the plasma formed on the screen surface contains metal nitrides formed by the combination of metal atoms expelled from the screen by nitrogen ion sputtering and excited nitrogen species. The nitrogen then diffuses into the material following decomposition of metal nitrides on the material surface, leading to nitriding. However, in ASPN, the nitriding rate is slower than that of DCPN because a layer of metal nitrides (deposit layer) is formed by the deposition of screen-derived metal nitrides on the material surface, which thickens over the course of the treatment time and inhibits nitrogen diffusion.⁴¹⁾ Our group has recently reported on DCPN in which a voltage is applied to both the material and the screen (S-DCPN).⁴¹⁾ In S-DCPN, the screen acts as a heater and reduces the cathodic voltage applied to the material, which mitigates the edge effect that occurs in DCPN. Furthermore, the nitriding rate was faster than that of DCPN and ASPN because of the removal of the deposit layer by sputtering on the material surface and the increased plasma formation area.

In ASPN and S-DCPN, steel materials are generally used as screen materials; there are few reports on the use of non-ferrous materials as screens. Our group reported plasma nitriding using chromium and titanium screens^42–45) and reported that there was limited nitrogen diffusion into the material in ASPN in this approach because the screen-derived chromium and titanium nitrides deposited on the material surface are highly stable and do not decompose easily. They also reported that the constituent atoms of the screens diffused into the material in both ASPN and S-DCPN. Other groups have reported on the use of Cr, Ti, Ni and solid carbon screens.^46–50) In one study, we applied ASPN and S-DCPN using a Ni screen and a steel screen to a low carbon steel (S15C), and evaluated how screen material affected the amount of nitrogen diffusion into the S15C steel.⁵¹⁾ We found that there was increased nitrogen diffusion into the S15C when using the Ni screen compared to the steel screen in both ASPN and S-DCPN, with greater differences observed for S-DCPN than ASPN. Therefore, it was theorized that performing S-DCPN using a Ni screen could increase the amount of nitrogen diffusion into the treated stainless steel and make the resulting S-phase and Sα-phase thicker than could be achieved using a steel screen.

To this end, the aim of this study was to thicken the S-phase and Sα-phase of stainless steel treated with low-temperature plasma nitriding using screen. Austenitic stainless steel (SUS304), ferritic stainless steel (SUS430), and duplex stainless steel (SUS329J4L) were treated with S-DCPN using a Ni screen or a steel plate cold commercial (SPCC) screen. The steel properties were compared after treatment using the different screen materials.

2. Experimental Methods

2.1 Sample materials

Austenitic stainless steel (SUS304), ferritic stainless steel (SUS430), and duplex stainless steel (SUS329J4L) were used in this study. The chemical compositions are listed in Table 1. Samples were cut from rods into a ϕ25 mm × 5 mm disc, following which the top surfaces were wet ground from #220 to #2000 and finally buffed to a mirror surface using alumina powder with a grain size of 1 µm.

Table 1 Chemical composition of SUS304, SUS430 and SUS329J4L (mass%).

2.2 Plasma nitriding treatment

A DC plasma-nitriding furnace (NDK Inc., model no. JIN-1S) was used for plasma nitriding. The screens were ϕ100 mm in diameter, 56 mm in height, and made of a Ni mesh (ϕ0.15 mm wire, 50 mesh, 49.7% open area ratio) or an expanded SPCC (LW = 4.0 mm, SW = 8.0 mm, T = 0.5 mm, W = 0.8 mm, 66.3% open area ratio). SUS430 rods were placed under the samples. The distance between the top of the samples and the top lid of the screens was 15 mm. Plasma nitriding was performed at 673 K for 300 min under a 75% N₂ + 25%₂ H₂ gas mixture atmosphere at a gas pressure of 100 Pa after the furnace was evacuated to less than 10 Pa.

2.3 Evaluation methods

To examine the plasma-nitrided sample surface, X-ray diffraction (XRD) measurements were performed on the plasma-nitrided samples using an X-ray diffractometer (RIGAKU, model no. RINT-2200). Cu-Kα radiation (wavelength λ = 0.15405 nm) was used as the X-ray source, and the tests were performed at a tube voltage of 40 kV and a tube current of 300 mA. For elemental analysis in the depth direction from the surface, glow discharge optical emission spectroscopy (GD-OES) was performed using a Marcus-type frequency glow discharge optical emission surface analyzer (Horiba, model no. GD-Profiler2). To analyze the cross-sectional and surface microstructures, electron probe microanalysis (EPMA) was performed using an electron probe microanalyzer (JEOL, model no. JXA-8230). To observe the cross-sectional microstructure, electrolytic corrosion was performed using 5 mass% oxalic acid as the corrosion solution (SUS304: 5 V, 15 s; SUS430: 3 V, 5 s; SUS329J4L: 5 V, 10 s). The hardness of the samples was measured using a micro Vickers hardness tester (Matsuzawa, model no. PMT-X7A). Polarization curves were generated to examine the pitting corrosion resistance in saltwater. The test areas of the samples were limited by covering them with Teflon tape with a ϕ6 mm hole. The test solution was a 3.5 mass% NaCl solution, the reference electrode was Ag/AgCl, and the counter electrode was Pt. A voltage between −1.0 V to +1.5 V was applied using a potentiostat (Hokuto Denko, model no. HA-501G) and the current density was recorded using a data logger (Graphtec, model no. GL200A-UM801). Following the corrosion test, the corroded areas of the samples were examined using an optical microscope (Olympus, model no. BX-60M).

3. Experimental Results

The XRD results for the nitrided samples are shown in Figs. 1(a)–(c) for SUS304, SUS430, and SUS329J4L, respectively. In all nitrided SUS304 samples, diffraction lines for the γ-phase (matrix) and expanded austenite (S) phase, were observed (Fig. 1(a)). The diffraction lines for the S phase were consistent with those of the γ-phase are wider and shifted to lower angles. In all nitrided SUS430 samples, the diffraction lines of the expanded ferrite (Sα) phase were detected and were consistent with those of the α-phase (matrix) shifted to lower angles (Fig. 1(b)). The CrN diffraction lines were also detected at approximately 63.3°, although this is not shown in Fig. 1(b) as these lines were low intensity. In all nitrided SUS329J4L samples, the diffraction lines of the γ-phase (matrix), α-phase (matrix), and S-phase were observed, but the diffraction lines of the Sα-phase were not observed (Fig. 1(c)). The shifts of the diffraction lines to lower angles and the widening of the diffraction lines for these expanded phases are caused by the supersaturation of N in the crystal lattices of the γ- and α-phases. This supersaturation leads to the introduction of compressive residual stress, an increase in stacking defect density, and large elastic strain on the crystal lattices; furthermore, it improves the mechanical properties of stainless steels, such as wear resistance and surface hardness.^3,12,17) Therefore, the expansion ratios of the lattice parameter were calculated from the diffraction angles of the expanded phase of each stainless steel sample as determined from the XRD results. The results are summarized in Table 2. In all steel grades, for the same lattice planes, the expansion rates when using the Ni screen are greater than those of the samples treated using the SPCC screen, which indicates that more N is solid-soluted in the crystal lattice. Comparisons of the expansion rates of the S phase in different lattice planes show that the expansion rate of the S(200) plane is greater than that of the S(111) plane, indicating that the crystal lattice expanded anisotropically. This was probably caused by the greater elastic modulus in the ⟨111⟩ direction than the ⟨200⟩ direction in the face centered cubic (fcc) structure of γ-Fe and the lower diffusion rate of N in the ⟨111⟩ direction compared to the ⟨200⟩ direction.^52–55)

Fig. 1

X-ray diffraction patterns of (a) SUS304, (b) SUS430 and (c) SUS329J4L samples treated by S-DCPN using Ni screen and SPCC screen.

Table 2 Expansion rate (%) of lattice parameter of SUS304, SUS430 and SUS329J4L samples treated by S-DCPN using Ni and SPCC screen.

Figures 2(a)–(c) show the profiles of the N element in SUS304, SUS430, and SUS329J4L by GD-OES. The use of the Ni screen increased nitrogen diffusion in all steel grades (Figs. 2(a)–(c)). This result is consistent with previous work on S-DCPN for S15C, in which the amount of nitrogen diffusion was greater when using the Ni screen than when using the SPCC screen.⁵¹⁾ This is because nickel nitrides from the Ni screen are less stable than iron nitrides from the SPCC screen and decompose more easily on the sample surface, resulting in more N diffusion into the sample. Comparing the N profiles of SUS304 and SUS329J4L (Figs. 2(a) and (c)) with that of SUS430 (Fig. 2(b)), the former had large N concentrations at 3–4 µm from the surface but limited diffusion depths, whereas the latter had limited N concentrations at 3–4 µm from the surfaces but large diffusion depth. This is thought to be due to the solid solubility of N atoms in α-Fe being less than that in γ-Fe,^6,56) and the diffusion rate of N atoms in α-Fe being greater than that in γ-Fe.^16,18,57)

Fig. 2

GD-OES nitrogen profiles of (a) SUS304, (b) SUS430 and (c) SUS329J4L samples treated by S-DCPN using Ni screen and SPCC screen.

The results of the cross-sectional microstructural observation by EPMA are shown in Figs. 3, 4, and 5 for SUS304, SUS430, and SUS329J4L, respectively. The inset figures in Fig. 4 show an enlarged view of the sample surface. The distinction between the α-phase and γ-phase in Fig. 5 is based on the results of the EPMA area analysis, where the Cr- and Mo-enriched phase is judged to be the α-phase and the Ni-enriched phase is judged to be the γ-phase. The thicknesses of the nitrided layers measured from Figs. 3, 5(a), and (b) are shown in Table 3. In all nitrided SUS304, nitrided layers were observed on the sample surfaces (Fig. 3 and Table 3); these layers were probably composed of the S-phase based on the S-phase identified by XRD (Fig. 1(a)). The S-phase of the sample treated using the Ni screen was approximately 1.7 times thicker than that of the sample treated using the SPCC screen. No clear nitrided layer was observed on the surfaces of any nitrided SUS430 samples (Fig. 4). This indicates that there is no significant difference in the microstructure between the Sα-phase and the α-phase (matrix).^13,19) Needle-like precipitates of iron nitrides were observed in the SUS430 samples, though they were not identified by XRD (Fig. 1(b)). Casteletti et al. found that the large stacking defect density and elastic strain on the crystal lattice in the expanded phase lead to wider asymmetric diffraction lines, making it difficult to identify phases with very small volume fractions, such as chromium nitrides and iron nitrides.¹²⁾ Therefore, this observation may be attributable to the formation of the Sα-phase. In addition, in the inset of Fig. 4, strongly etched areas are visible on the top surfaces of the samples. As shown in the XRD results, although the intensity of the diffraction lines was low, CrN was identified, which may be attributed to the decrease in corrosion resistance due to CrN precipitation in these areas. In all nitrided SUS329J4L samples, nitrided layers were formed independent of the α-phase or γ-phase. Hence, elemental analysis by EPMA was conducted at the center of each nitrided layer; the detected N concentrations are listed in Table 4. The N concentrations detected in the center of the nitride layers formed in the α- and γ-phases were similar. In combination with the fact that no Sα-phase was identified by XRD unlike the S-phase (Fig. 1(c)), this result suggests that the S-phase was formed in both the γ-phase and the α-phase. Similar results have been reported in low-temperature nitriding of duplex stainless steels and a phase transformation from the α-phase to the S-phase has been proposed.^{13,14,17,18,21,22,57)} It is possible that N (a strong γ-stabilizing element) diffused into the α-phase and caused the phase transformation, as shown in reaction (1), and the S-phase was formed thereafter.   

\begin{equation} \alpha \to (\text{S}\alpha,\gamma)\to \text{S} \end{equation} (1)

In addition, SUS329J4L samples exhibited needle-like structures in the S-phase that formed in the α-phase, consistent with previous reports on low-temperature nitriding of duplex stainless steel. Pinedo et al. reported that these structures were the deformation zone formed at the interface between the S-phase and the matrix due to the difference in crystal expansion between the S-phase and the ferrite structure of the matrix,¹⁶⁾ while Tschiptschin et al. reported that although iron nitrides were identified by XRD, they could not distinguish whether they arose due to the deformation zone or precipitation of iron nitrides.⁵⁸⁾ As shown in Table 3, the S-phase formed on the sample treated using the Ni screen was approximately 1.3 times thicker than that treated using the SPCC screen. Furthermore, comparisons of the thickness of the S-phase formed in the α- and γ-phases when using the same screen material showed that the S-phases formed in the α-phase were thicker than those in the γ-phase. This is because N atoms preferentially diffuse into the α-phase, due to its increased diffusion coefficient relative to the γ-phase.^6,16,18,57) The low-magnification images shown in Fig. 5(c) and (d) show that the S-phase is locally thin in the sample treated using the SPCC screen without distinguishing between the α-phase and γ-phase, whereas no such region was observed in the sample treated using the Ni screen.

Fig. 3

Cross-sectional microstructure of SUS304 samples treated by S-DCPN using (a) Ni screen and (b) SPCC screen.

Fig. 4

Cross-sectional microstructure of SUS430 samples treated by S-DCPN using (a) Ni screen and (b) SPCC screen.

Fig. 5

Cross-sectional microstructure of SUS329J4L samples treated by S-DCPN using (a), (c) Ni screen and (b), (d) SPCC screen.

Table 3 Thickness (µm) of nitrided layer of SUS304 and SUS329J4L samples treated by S-DCPN using Ni and SPCC screen.

Table 4 N concentration (at%) at center of nitrided layer of SUS329J4L samples treated by S-DCPN using Ni and SPCC screen.

The results of the surface hardness test are presented in Fig. 6, which shows that the surface hardness of all treated steels was higher than that of the untreated samples. The increase in surface hardness was attributed to the formation of the S-phase in SUS304 and SUS329J4L and the formation of the Sα-phase and precipitation strengthening with iron nitrides and CrN in SUS430. In addition, the surface hardness of the samples treated with the Ni screen was greater than that of the samples treated with the SPCC screen. In SUS329J4L, the surface hardness was comparable between the α-phase and γ-phase for the samples treated with the Ni screen, while areas of different surface hardness were observed between the phases for samples treated with the SPCC screen. Therefore, surface microstructure observations (Fig. 7) and elemental analysis by EPMA (Table 5) were performed to examine the surface microstructures after the Vickers hardness tests were completed. In SUS329J4L, N concentrations in the α-phase and γ-phase were similar in the sample treated using Ni screen; conversely, the N concentrations differed by approximately 30 at% in the high surface hardness region and approximately 20 at% in the low surface hardness region in the sample treated using the SPCC screen (Fig. 7 and Table 5). This suggests that the formation of regions with different surface hardness is due to the non-uniform distribution of N in the sample.

Fig. 6

Surface hardness of SUS304, SUS430 and SUS329J4L samples treated by S-DCPN using Ni screen and SPCC screen.

Fig. 7

Surface microstructure of Vickers hardness tested areas of SUS329J4L samples treated by S-DCPN using (a) Ni screen and (b) SPCC screen.

Table 5 Elemental analysis in Fig. 7(a) and (b) (at%).

In the polarization test conducted to examine the pitting corrosion resistance to saltwater, the potentials at the current density i = 1 A m⁻² were defined as the pitting corrosion potential, as shown in Fig. 8. Post-polarization test samples are shown in Fig. 9. It is worth noting that in SUS329J4L, the evaluation was restricted to the gap areas between the Teflon tape and the examination surface because of its high resistance to pitting corrosion, and only the corrosion marks were examined. Several theories have been proposed to explain the increased pitting corrosion resistance of stainless steels resulting from the formation of the S-phase and Sα-phase. The most popular theory proposes that solid solution N in these phases mitigates the change in pH in the pit.^2,8,13,22) When the passive film of stainless steel is destroyed by Cl⁻, reaction (2) occurs in the localized pits, which suppresses the formation of a local environment with low pH and high Cl⁻ concentration that causes active dissolution and in turn suppresses pit growth and promotes re-passivation.   

\begin{equation} [\text{N}] + \text{4H$^{+}$} + \text{3e$^{-}$}\to \text{NH$_{4}{}^{+}$} \end{equation} (2)

As shown in Figs. 8 and 9, in all nitrided SUS304 samples, the pitting corrosion potentials increased and the corroded areas decreased compared to the untreated sample, indicating improved pitting corrosion resistance. In addition, the pitting potential of the sample treated using the Ni screen was greater than that of the sample treated using the SPCC screen, indicating better resistance to pitting corrosion. In all nitrided SUS430 samples, the pitting corrosion potentials decreased and the corroded areas increased compared to the untreated sample, indicating reduced pitting corrosion resistance. This is probably due to the decrease in solid solution Cr resulting from CrN precipitation, as suggested by the XRD results (Fig. 1(b)) and cross-sectional microstructure observation (Fig. 4). The pitting corrosion potential of the sample treated using the Ni screen was less than that of the sample treated using the SPCC screen, indicating that the pitting corrosion resistance was further decreased by treatment with a Ni screen. In all nitrided SUS329J4L samples, the pitting corrosion potentials increased and the corroded areas decreased compared to the untreated sample, indicating improved pitting corrosion resistance. In addition, the pitting potential of the sample treated with the Ni screen was slightly greater than that of the sample treated with the SPCC screen. As shown in Fig. 8, SUS329J4L has excellent pitting corrosion resistance, unlike SUS304 and SUS430; therefore, there was no significant difference in the pitting corrosion potential between the samples treated using the Ni screen and the SPCC screen in this study. In order to further analyze the crevice corrosion areas of SUS329J4L, elemental analysis by EPMA was performed. The crevice corrosion areas of the untreated sample and the sample treated with the Ni screen are shown in Fig. 10. The crevice corrosion area of the sample treated using the SPCC screen had comparable morphology to the sample treated using the Ni screen (data not shown). As shown in Fig. 10, extensive corrosion marks were observed in the untreated sample and the α-phase, in which more Mo was detected, was more corroded than the γ-phase, in which more Ni was detected. This indicates that the corrosion progress differed depending on the phase. In the nitrided sample, the corrosion marks were narrower and deeper than those in the untreated sample; the variation in corrosion progress between phases was not observed. It is proposed that the surface pitting corrosion resistance became uniform between phases following nitriding due to the formation of an S-phase that had the same N concentration as the α- and γ-phases, as shown in the results of the elemental analysis by EPMA (Table 5).

Fig. 8

Pitting corrosion potential of SUS304, SUS430 and SUS329J4L samples treated by S-DCPN using Ni and SPCC screen.

Fig. 9

Surface microstructure of corrosion tested areas in the 3.5 mass% NaCl of SUS304, SUS430 and SUS329J4L samples treated by S-DCPN using Ni screen and SPCC screen.

Fig. 10

Elemental analysis of crevice corrosion areas in the 3.5 mass% NaCl of SUS329J4L untreated sample and SUS329J4L sample treated by S-DCPN using Ni screen.

4. Discussion

4.1 Effect of using a Ni screen for low-temperature plasma nitriding of austenitic stainless steels

As shown in the GD-OES results (Fig. 2(a)), nitrogen diffusion was increased when using a Ni screen. Therefore, as shown in the XRD results (Fig. 1(a) and Table 2) and cross-sectional microstructure observations (Fig. 3 and Table 3), a thicker S phase with more solid solution N in the crystal lattice was formed. Accordingly, as shown in the results of the surface hardness test (Fig. 6) and pitting corrosion tests (Figs. 8 and 9), the surface hardness and pitting corrosion resistance of the samples increased.

4.2 Effect of using a Ni screen for low-temperature plasma nitriding of ferritic stainless steels

As shown in the GD-OES results (Fig. 2(b)), nitrogen diffusion was increased by using the Ni screen. Therefore, as shown in the XRD results (Fig. 1(b) and Table 2) and the surface hardness test (Fig. 6), an Sα-phase with more solid solution N in the crystal lattice was formed, and the surface hardness was increased by the enhanced nitride precipitation and formation of the Sα-phase. However, CrN precipitation was also accelerated, which may have decreased the pitting corrosion resistance, as shown in the results of the pitting corrosion tests (Figs. 8 and 9). These nitride precipitates are not observed in SUS304 and SUS329J4L. This may result from the reduced solubility limit of N in the Sα-phase compared to the S-phase. This is consistent with previous work indicating that the solid solubility limit of N in α-Fe is smaller than that in γ-Fe^6,56) and the lattice parameter expansion rate of the Sα-phase is smaller than that of the S phase, as shown in the XRD results (Table 2). In this study, CrN and iron nitrides may have precipitated because of N diffusion that exceeded the solid solubility limit of the Sα-phase.

4.3 Effect of using a Ni screen on low-temperature plasma nitriding of duplex stainless steels

As shown in the GD-OES results (Fig. 2(c)), nitrogen diffusion was increased by using the Ni screen. Therefore, as shown in the XRD results (Fig. 1(c) and Table 2) and cross-sectional microstructure observations (Fig. 5 and Table 3), a thicker S phase with more solid solution N in the crystal lattice was formed. Accordingly, as shown in the results of the surface hardness test (Fig. 6) and pitting corrosion test (Figs. 8 and 9), the surface hardness and pitting corrosion resistance of the samples increased. In the sample treated using the SPCC screen, cross-sectional microstructural observations (Fig. 5(c) and (d)) and surface hardness tests (Fig. 6, Fig. 7, and Table 5) revealed the formation of a localized thin S-phase as well as phases with different hardness and N concentrations at the sample surface, suggesting a non-uniform N distribution. This may be due to the varying diffusivities of N in different lattice planes and the fact that iron nitrides from the SPCC screen decomposed less than nickel nitrides from the Ni screen. Gallo and Dong reported that the diffusivity of N was larger in the (100) plane than in the (111) plane,⁵⁹⁾ and Martinavicius et al. reported that the diffusion coefficient of N in (100)-oriented samples is approximately twice that in (111)-oriented samples in austenitic stainless steels.⁶⁰⁾ Furthermore, Borgioli et al. and Mändl and Rauschenbach reported that the magnitude of the lattice expansion rate is as follows: (200) > (311) > (111) and (222).^3,53) We have also reported that nickel nitrides decompose more easily on the sample surface than iron nitrides; thus, more N diffuses into the sample.⁵¹⁾ In this study, as shown in the GD-OES results (Fig. 2(c)), there was increased nitrogen diffusion in the samples treated using the Ni screen compared to those treated using the SPCC screen. These suggest that iron nitrides decompose easily on the lattice planes with greater N diffusibility and less easily on the lattice planes with small N diffusibility, resulting in the formation of a thin S phase. This effect also results in phases with different surface hardness and N concentrations. Furthermore, the uniform distribution of N in the sample treated using the Ni screen may be attributable to the N diffusion being similar in both the small diffusibility and large diffusibility lattice planes of the sample surface. This is because of the easy decomposition of nickel nitrides, which results in comparable N diffusion in all lattice planes.

5. Conclusion

In this study, we aimed to thicken the S-phase and Sα-phase of stainless steel treated with low-temperature plasma nitriding using a screen. Austenitic stainless steel (SUS304), ferritic stainless steel (SUS430), and duplex stainless steel (SUS329J4L) were subjected to S-DCPN using a Ni screen and a steel plate cold commercial (SPCC) screen, and we compared the differences in the steel samples between screen materials. The following conclusions were drawn:

1. (1)    For S-DCPN treatment of SUS304, the S-phase became thicker when using the Ni screen than when using the SPCC screen because of the increase in nitrogen diffusion. Accordingly, the surface hardness and pitting corrosion resistance of SUS304 were increased following treatment with a Ni screen compared to an SPCC screen.
2. (2)    For S-DCPN treatment of SUS430, there was increased nitrogen diffusion when using a Ni screen compared to an SPCC screen, but the thickness of the resulting Sα-phase could not be measured. Although the surface hardness was increased, the pitting corrosion resistance was lower than that of the untreated sample owing to CrN precipitation using both screens and was lower with the Ni screen than the SPCC screen.
3. (3)    For S-DCPN treatment of SUS329J4L, the S-phase was thicker for samples treated using the Ni screen than those treated using the SPCC screen because the amount of nitrogen diffusion increased. Accordingly, the surface hardness and pitting corrosion resistance increased in the sample treated using the Ni screen-treated sample relative to the sample treated using the SPCC screen.
4. (4)    In low-temperature plasma nitriding using a screen, the use of a Ni screen results in a thicker S-phase in austenitic and duplex stainless steels due to an increase in the amount of nitrogen diffusion. This in turn increases the surface hardness and pitting resistance compared to the use of a conventional steel screen. In ferritic stainless steels, the use of a Ni screen results in an increase in the amount of nitrogen diffusion and surface hardness compared to the conventional steel screen. However, it is necessary to account for the decrease in pitting corrosion resistance due to the accompanying precipitation of CrN.

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