Intergranular and pitting corrosion susceptibility of sensitized selective laser melted 316L stainless steel

Abstract

The intergranular corrosion (IGC) and pitting corrosion behavior of selective laser melted (SLMed) 316L stainless steel (316L SS) after sensitization were systematically investigated and compared with its wrought counterpart. The results reveal that SLMed 316L SS exhibits inferior IGC resistance compared with the wrought material due to its lower proportion of Σ3 coincident-site lattice (CSL) boundaries. However, it shows significantly higher pitting corrosion resistance, which is attributed to the refinement of non-metallic inclusions. Following sensitization treatment, both materials experienced a significant reduction in pitting resistance alongside an increase in metastable pitting events. The underlying degradation mechanisms are elucidated through the analysis of preferential initiation sites and passive film composition evolution.

Keywords

selective laser melted stainless steel intergranular corrosion pitting corrosion passive film

Introduction

316L stainless steel (316L SS) is widely used as a high-temperature structural material in nuclear reactors, steam-generating plants, and chemical processing due to its excellent corrosion resistance and high-temperature mechanical properties.^1–3 However, under harsh environmental conditions (e.g., prolonged exposure to temperatures between 500 and 850 °C), conventional SSs often undergo sensitization due to chromium carbide precipitation along grain boundaries (GBs), making the material susceptible to intergranular corrosion (IGC).^4–6 Selective laser melting (SLM), also known as laser powder bed fusion (LPBF), is an emerging technology that fabricates near-net-shape parts using high-energy laser beams. Compared with traditional manufacturing technologies such as casting, rolling, and forging, SLM is both technologically advantageous and economically competitive, as it enables the rapid fabrication of complex components without additional costs.^7–9 Therefore, SLM technology has attracted considerable attention for manufacturing high-temperature structural components and has been proven as an ideal technique for 316L SS.^7,10

The rapid melting and solidification, coupled with cyclic heating and cooling during the deposition of subsequent layers in SLM manufacturing, result in microstructural differences between selective laser melted (SLMed) 316L SS and conventional 316L SS.^11,12 Numerous reports have focused on the mechanical properties and corrosion behavior of SLMed SSs.^7,10^13–16 Previous reports suggest that SLMed 316L SS exhibits superior mechanical properties and corrosion resistance compared to its conventional counterpart.^15,16 However, only a few studies have investigated the sensitization behavior of SLMed 316L SS, and the reported findings are inconsistent. Man et al.¹⁷ found that SLMed 316L SS could be sensitized at 650 °C when the treatment duration exceeded 6 h. Macatangay et al.¹⁸ reported that SLMed 316L SS was more susceptible to sensitization than its wrought counterpart at 675 °C. The SLMed 316L SS was sensitized after only 1 h of treatment at 675 °C, whereas the wrought 316L SS required a much longer time (∼24 h) at the same temperature. However, Laleh et al.¹⁹ found that the IGC resistance of SLMed 316L SS was higher than that of conventional 316L SS, attributing this to the absence of Cr-rich precipitates in SLMed specimens after long-term sensitization heat treatment. In contrast to the above findings, Snitzer et al.²⁰ and Macatangay et al.²¹ reported that neither SLMed nor wrought 316L SS exhibited sensitization after 24 h at 700 °C. Therefore, the extent and mechanisms of IGC in SLMed 316L SS still require further investigation.

Sensitization induces chromium depletion in regions adjacent to GBs, resulting in decreased corrosion resistance of SS. While numerous studies have investigated the influence of sensitization on pitting corrosion in conventional SS,^5,6,22–25 this topic has not been adequately considered in SLMed SS. Cheng et al.⁶ demonstrated that sensitization reduced the pitting corrosion resistance of 304 SS due to the formation of Cr-rich precipitates. The detrimental effect of sensitization on pitting resistance and passive film stability was also confirmed by Lv et al.²⁴ Ida et al.²⁵ investigated pitting initiation in sensitized 304 stainless steel using micro-scale localized polarization and microscopic characterization. Their study revealed that pitting corrosion did not occur at sensitized GBs without sulfide inclusions, demonstrating that MnS inclusions are necessary for pit initiation even at sensitized GBs. However, the additive manufacturing process and resultant solidification-induced microstructural features unique to SLMed 316L SS — such as porosity, residual stress, high dislocation density, and limited twin boundaries — significantly influence its corrosion performance.^3,26,27 Nevertheless, the role of sensitization in pitting corrosion of SLMed SS remains poorly understood. A comprehensive understanding of sensitization effects on the corrosion behavior of SLMed 316L SS would be beneficial for optimizing post-processing parameters in high-temperature engineering applications.

In this study, the IGC susceptibility and pitting corrosion behavior of sensitized SLMed 316L SS were investigated and systematically compared with those of wrought counterparts. The following characterization techniques were employed: Electron backscatter diffraction (EBSD), scanning electron microscopy (SEM) equipped with an Oxford AZtec energy dispersive spectroscopy (EDS) system, different electrochemical measurements, and X-ray photoelectron spectroscopy (XPS).

Experimental methods

Materials and heat treatment

Commercially available 316L SS powder with a particle size distribution of 15∼53 μm was used to fabricate SLMed samples. A P400 SLM machine was employed to produce cubic specimens with dimensions of 1.0 × 1.0 × 1.0 cm³. Prior to SLM fabrication, the build plate was pre-heated to 150 °C and the build chamber was purged with purified argon until the oxygen level decreased below 100 ppm. The main processing parameters included a laser power of 150 W, a scanning speed of 600 mm/s, a hatch spacing of 80 μm, and a layer thickness of 35 μm. The scanning pattern was bidirectional with 67° angle between successive layers. The compactness of the SLMed sample was approximately 99.7%, corresponding to a porosity of less than 0.03%. Electrochemical measurements and microstructural analyses were conducted on the surface perpendicular to the build direction. For comparison, a wrought 316L SS plate was also used in this study. Prior to sensitization, the wrought material was solution annealed at 1100 °C for 1 h. Table 1 lists the chemical composition of as-built SLMed and wrought 316L SSs, indicating that the two materials have comparable chemical contents. For the sensitization treatment, the samples were heat-treated at 650 °C for 40 h, followed by immediate water quenching. This specific condition was selected to induce a substantial degree of sensitization (DOS), as it lies within the well-established critical temperature-time regime for the precipitation of Cr-rich carbides in 316L SS, as extensively documented in the literature.^19,28

Table 1.

The composition of as-built SLMed and wrought 316L SSs.

Element (wt.%)	C	Cr	Ni	Mn	Si	P	S	Mo	Fe
as-built SLMed	0.02	17.52	9.54	1.04	1.15	0.02	0.01	1.87	Bal.
wrought	0.02	17.01	10.85	0.98	0.76	0.01	0.01	1.14	Bal.

Electrochemical experiments

Electrochemical measurements were performed in CS350 by a conventional three-electrode cell system. A saturated calomel electrode (SCE) and a platinum plate were employed as the reference electrode and counter electrode, respectively. Samples for electrochemical tests were cold-mounted in epoxy resin, with an exposed surface area of approximately 1 cm² serving as the working electrode. Prior to each electrochemical measurement, the sample surface was sequentially ground with silicon carbide abrasive paper (from 180 to 2000 grit), rinsed with ethanol, and dried. Before the electrochemical measurements, the open-circuit potential (OCP) was monitored for 30 min to ensure it stabilized. All electrochemical tests were carried out at room temperature and repeated five times to ensure the reproducibility and consistency of the results.

The DOS was evaluated using double-loop electrochemical potentiokinetic reaction (DL-EPR) in 0.5 M H₂SO₄+ 0.01 M KSCN solution. The potential was first scanned in the anodic direction from the OCP to +0.3 V_SCE, and then immediately reversed back to the OCP at a scan rate of 100 mV/min.

To analyze the effect of sensitization on the pitting corrosion behavior of SLMed 316L SS under controlled and reproducible conditions, all pitting-related electrochemical tests were conducted in a neutral 3.5 wt.% NaCl solution at room temperature. Although this benchmark electrolyte does not replicate the full complexity of industrial environments (e.g., with variations in temperature, pH, or the presence of SO₄²⁻ ions), it allows for the isolation of microstructural effects on pitting behavior and passive film properties. Potentiodynamic polarization tests were conducted by scanning from 0.3 V below the OCP at a rate of 0.5 mV/s until the anodic current density exceeded 1 mA/cm². To study metastable pitting behavior, potentiostatic polarization was carried out in the same solution at a potential of 0.15 V_SCE below the average pitting potential (E_pit) of the respective sample, as determined from the potentiodynamic polarization tests. During these tests, all samples (with an exposed surface area of 0.1 cm²) were monitored at a frequency of 20 Hz for 1 h to minimize the overlap of current transients.

Electrochemical impedance spectra (EIS) and Mott-Schottky measurements were conducted at the OCP after each sample was immersed in 3.5 wt.% NaCl solution for 2 h. EIS measurements were performed over a frequency range from 100 kHz to 10 mHz with a signal amplitude perturbation voltage of 10 mV. The experimental data were analyzed by Zview software. Mott-Schottky measurements were carried out by scanning the potential in the cathodic direction from the film formation potential down to −1.5 V_SCE with a scan rate of 40 mV/steps. A fixed frequency of 1000 Hz and a perturbation voltage of 10 mV were used. According to the Mott-Schottky theory, the capacitance of the Helmholtz layer (C_H) can be neglected. Thus, the semiconductor properties of the passive film can be described as follows:²⁹

n - type : \frac{1}{C^{2}} \approx \frac{1}{C_{S C}^{2}} = \frac{2}{ε ε_{0} e N_{D}} (E - E_{F B} - \frac{k T}{q})

(1)

p - type : \frac{1}{C^{2}} \approx \frac{1}{C_{S C}^{2}} = - \frac{2}{ε ε_{0} e N_{A}} (E - E_{F B} - \frac{k T}{q})

(2)

where C_SC is the space charge capacitance, ε is the relative dielectric constant of the passive film (which is assumed as 12 for stainless steel), ε₀ is the vacuum permittivity (8.854 × 10⁻¹⁴ Fm⁻¹), e is the electron charge, k is the Boltzmann's constant, E_FB is the flat band potential. N_D and N_A are the donor and acceptor concentrations, respectively. The values of N_D and N_A could be determined by the positive and negative slopes of the linear region in Mott-Schottky curves.

Microstructure analyses

EBSD measurements were employed to analyze the microstructures of SLMed and wrought 316L SSs. The measurements were performed using an FEI Quanta 650 SEM equipped with an Oxford Instruments EBSD detector. The acceleration voltage and scanning step size were set to 20 kV and 0.8 μm, respectively. Prior to EBSD analysis, all samples were mechanically polished with a non-crystallizing colloidal silica suspension to remove surface deformation layers induced by preliminary diamond abrasive polishing. The acquired EBSD data were processed using HKL Channel 5 software to determine grain orientation, GB character, and dislocation distributions.

Statistical analysis of the macroscopic distribution, chemical composition and quantity of inclusions was carried out using a Zeiss Sigma 300 SEM equipped with an Oxford-Aztec EDS system, operating at an acceleration voltage of 15 kV. The analysis was conducted on polished surfaces over a scanned area of 5 × 5 mm².

Following DL-EPR measurements, the samples were examined by optical microscopy to evaluate intergranular attack. To identify metastable pitting initiation sites, both the morphology and chemical composition of pits formed during potentiostatic polarization were analyzed using SEM combined with EDS.

XPS measurements were conducted to analyze the chemical composition of the passive films. Prior to XPS analysis, samples were immersed in 3.5 wt.% NaCl solution for 2 h to form stable passive film. The measurements were conducted using a Thermo ESCALAB 250Xi XPS spectrometer without any sputtering technique. The X-ray source employed was a monochromator Al Ka (1486.6 eV, 15 kV, 14.9 mA). The spectrometer was operated with a constant pass energy of 20 eV. XPS data were analyzed using commercial XPSpeak (version 4.1) software with Shirley background subtraction. The C1 s peak from adventitious carbon at 284.8 eV was used as calibration peak to correct charging shifts, and all the other peaks were calibrated using the C1 s peak.

Results

Figure 1 shows the inverse pole figure (IPF) maps, GB maps and kernel average misorientation (KAM) maps of SLMed and wrought 316L SSs analyzed by EBSD. As shown in Figures 1(a1) and (b1), the two materials exhibit distinct differences in grain morphology and GB characteristics. The SLMed 316L SS displays wavelike GBs, whereas the wrought 316L SS consists mainly of grains with nearly straight GBs. Based on EBSD data, grain size and GB length of the two 316L SSs were statistically analyzed. The average grain size of SLMed and wrought 316L SSs was evaluated to be 10.05 μm and 9.97 μm, respectively, indicating comparable grain dimensions. However, the total GB length of the SLMed 316L SS (18.92 mm) was approximately 2.18 times that of the wrought sample (8.66 mm). As shown in Figures 1(a2) and (b2), different types of GBs in the maps are presented with distinct colors, with black and red lines denoting high-angle GBs (HAGBs, >15^o of misorientation) and ∑3 coincident-site lattice (CSL) boundaries, respectively. Notably, the SLMed 316L SS exhibits a significantly lower proportion of ∑3 CSL boundaries (9.71%) than the wrought sample (18.90%). The KAM maps can be used to characterize the dislocation distribution in the two steels.^30,31 As shown in Figures 1(a3) and (b3), dislocations in the SLMed sample are more concentrated near GBs compared to those in the wrought sample. In addition, the mean KAM value of the SLMed 316L SS (∼0.65) is significantly higher than that of the wrought sample (∼0.26), indicating a greater dislocation density in the former. Therefore, the EBSD analysis reveals distinct differences in grain morphology, GB characteristics, and dislocation concentration between the SLMed and wrought samples. These microstructural variations are expected to significantly influence their IGC resistance, pitting corrosion behavior, and passive film stability.

Figure 1.

The EBSD results of (a1-a3) SLMed and (b1-b3) wrought 316L SSs: (a1) and (b1) IPF maps; (a2) and (b2) GB maps; (a3) and (b3) KAM maps.

Figure 2 displays typical SEM micrographs of inclusions observed in the scanning areas of SLMed and wrought 316L SSs. Based on their morphological and compositional contrasts under SEM-EDS mapping, three distinct inclusion types were identified in both materials: sulfide inclusions (Figures 2(a1) and (b1)), oxide inclusions (Figures 2(a2) and (b2)), and composite sulfide-oxide inclusions (Figures 2(a3) and (b3)). The quantitative statistics of inclusion characteristics — including number density, average area, and equivalent circle diameter (ECD) — are summarized in Tables 2 and 3. In the SLMed 316L SS, the number densities of sulfide, oxide, and composite sulfide-oxide inclusions are 0.72, 17.76 and 17.60 mm⁻², respectively, with corresponding average areas of 4.06, 6.35 and 4.68 μm². In contrast, the wrought sample exhibits higher number densities and larger average areas for all three inclusion types compared to the SLMed material. For example, the composite sulfide-oxide inclusions in the wrought sample show a number density of 19.32 mm⁻² and an average area of 5.60 um². Moreover, the average ECD of all three inclusion types in the SLMed sample is smaller than that in the wrought sample (Tables 2 and 3). In summary, the quantitative analysis confirms that the SLMed 316L SS possesses a significantly refined and less dense inclusion population compared to the wrought material.

Figure 2.

The typical SEM micrographs of (a1, b1) sulfide inclusions, (a2, b2) oxide inclusions, (a3, b3) composite sulfide-oxide inclusions in (a1-a3) SLMed and (b1-b3) wrought 316L SSs within the scanning area.

Table 2.

The quantitative statistical results of inclusion characteristics in as-built SLMed 316L SS.

As-built SLMed	Number density (mm⁻²)	Average area (μm²)	ECD (μm)
As-built SLMed	Number density (mm⁻²)	Average area (μm²)	Max.	Min.	Average
Sulfide inclusions	0.72	4.06	3.47	0.61	2.16
Oxide inclusions	17.76	6.35	8.08	0.47	2.83
Composite sulfide-oxide inclusions	17.60	4.68	10.50	0.58	2.41

Table 3.

The quantitative statistical results of inclusion characteristics in wrought 316L SS.

Wrought	Number density (mm⁻²)	Average area (μm²)	ECD (μm)
Wrought	Number density (mm⁻²)	Average area (μm²)	Max.	Min.	Average
Sulfide inclusions	2.48	5.16	4.62	0.74	2.53
Oxide inclusions	20.24	7.52	20.61	0.52	3.08
Composite sulfide-oxide inclusions	19.32	5.60	14.84	0.68	2.64

Figure 3 shows the DL-EPR curves and the corresponding corroded surface microstructures of SLMed and wrought 316L SSs before and after sensitization. Distinct reactivation peaks are observed during the reverse scans for both sensitized SLMed (Figure 3(a)) and sensitized wrought (Figure 3(b)) samples. The DOS value was calculated from the ratio of the reverse scan reactivation peak current density (i_r) to the forward scan activation peak current density (i_a).²² After sensitization, both 316L SSs show increased DOS values, with the SLMed sample exhibiting a significantly higher DOS value (6.26%) than the wrought sample (2.15%). This result indicates that the SLMed 316L SS is more susceptible to IGC than its wrought counterpart.

Figure 3.

DL-EPR curves and surface corroded microstructures for SLMed and wrought 316L SSs: DL-EPR curves for (a) SLMed and (b) wrought 316L SSs before and after sensitization; morphology of (c, d) SLMed 316L SS, (e, f) wrought 316L SS after the DL-EPR test.

After DL-EPR measurements, the microstructure of the as-built SLMed 316L SS displays melt pool boundary attack, with some re-melted grains visible within the melt pool (Figure 3(c)). The sensitized SLMed sample shows localized melt pool boundary attack accompanied by distinct GB etching, suggesting the presence of chromium-depleted zones. Additionally, severe corrosion is also found in refined grains, as indicated by the arrows in Figure 3(d). In contrast, the as-received wrought sample shows a step-like morphology (Figure 3(e)), while the sensitized wrought sample develops a characteristic intergranular ditch structure (Figure 3(f)). The corrosion morphologies confirm that the SLMed sample suffers more severe intergranular attack than the wrought sample, which is consistent with the calculated DOS values.

Figure 4 presents the typical potentiodynamic polarization curves and the variation of E_pit for SLMed and wrought 316L SSs before and after sensitization. The sensitized samples exhibit lower corrosion potentials than their as-built SLMed and as-received wrought counterparts, suggesting reduced general corrosion resistance in NaCl solution after sensitization. A wide passive region is observed in all samples, indicating typical passive behavior in this environment. Sensitization increases the passive current density and promotes the occurrence of metastable pitting prior to stable pitting in both materials. Notably, the SLMed 316L SS shows a lower passive current density than the wrought sample, consistent with the findings reported by Man et al.¹⁵ Previous studies have indicated that the passive current density approximately correlates with the dissolution rate of the passive film during the passive process.³² Therefore, the lower passive current density of the SLMed sample suggests a lower dissolution rate of the passive film, implying that the passive film formed on the SLMed sample possesses superior protective ability compared to that on the wrought sample. However, this protective ability deteriorates after sensitization. As shown in Figure 4(c), the E_pit of SLMed 316L SS decreases from 0.57 V_SCE to 0.45 V_SCE after sensitization, while that of the wrought sample decreases from 0.45 V_SCE to 0.24 V_SCE. These results indicate that although sensitization undermines the pitting corrosion resistance of both materials, the SLMed 316L SS demonstrates superior retention of its protective properties compared to the wrought counterpart.

Figure 4.

Typical potentiodynamic polarization curves: (a) SLMed and (b) wrought 316L SSs before and after sensitization, (c) E_pit for all samples.

Figure 5 shows the metastable pitting current transients recorded at an applied potential of 0.15 V_SCE below E_pit in 3.5 wt% NaCl solution. These current transients reflect the initiation, early growth, and repassivation of metastable pits.³³ All samples exhibit similar trends in current fluctuations during the potentiostatic polarization process. The background current decays sharply within the first 500 s, then gradually decreases over time. The as-built SLMed 316L SS shows a higher frequency of current fluctuations within the first 500 s, followed by fewer current fluctuation spikes in the final stage of the potentiostatic polarization test. In contrast, the wrought 316L SS displays sustained current fluctuations throughout the entire test period. The number of current transients in the wrought sample is significantly greater than that in the SLMed material, indicating lower metastable pitting activity in the SLMed material. Both the magnitude and frequency of the metastable pitting current transients increase after sensitization, demonstrating that sensitization accelerates the dissolution kinetics of metastable pits.

Figure 5.

Current transients of SLMed and wrought 316L SSs under potentiostatic polarization in 3.5 wt.% NaCl solution at the applied potential.

The cumulative charge passed through metastable pits in each sample (Q_pit, in μC) was obtained by integrating all current peaks from rise to recovery during the entire metastable pitting process, which corresponds to the amount of metal dissolved within the pits.³⁴ Therefore, Q_pit reflects the pitting severity of each sample. The metastable pit initiation frequency (λ) was calculated by counting the number of events within 600 s intervals, then dividing by the interval duration and the electrode surface area (0.1 cm²). Note that all events with currents greater than 5 nA were considered metastable pits. The pit initiation frequency was plotted at the midpoint of each interval, and the results are shown in Figure 6. The wide range of cumulative charge values and the large scatter in pitting frequency demonstrate that metastable pitting is inherently random and stochastic. As shown in Figure 6, the Q_pit values for the as-built SLMed and as-received wrought 316L SSs are 0.80 μC and 6.56 μC, respectively. After sensitization, the Q_pit increases for both samples, indicating a rise in metastable pitting susceptibility. The wrought sample exhibits significantly higher metastable pitting susceptibility than the SLMed sample, as metastable pits with greater severity (higher Q_pit) are more likely to evolve into stable pits. As shown in Figure 6(b), the metastable pit initiation frequency (λ) decreases over time due to the consumption of a limited number of available initiation sites on all sample surfaces.³⁵ Similarly, the higher metastable pitting initiation frequency observed in the sensitized SLMed and wrought samples indicates an increased probability of pit formation, further confirming the degradation of pitting resistance after sensitization. Moreover, the as-built SLMed sample exhibits the lowest cumulative charge (Q_pit) and metastable pit frequency (λ), indicating superior resistance to metastable pitting. Crucially, even after sensitization, the SLMed sample suffers far less damage than the sensitized wrought sample. These statistical results demonstrate that sensitization impairs the pitting resistance of both SLMed and wrought samples. Nevertheless, the SLMed sample maintains significantly higher pitting resistance than the wrought sample, aligning with the potentiodynamic polarization results.

Figure 6.

(a) The cumulative charge passed through metastable pits and (b) metastable pit initiation frequency for SLMed and wrought 316L SSs during potentiostatic polarization at applied potential in 3.5 wt.% NaCl solution.

Figures 7 and 8 show the typical SEM morphologies and corresponding EDS spectra of pit initiation sites in SLMed and wrought 316L SSs after potentiostatic polarization at 0.15 V_SCE below their respective E_pit in 3.5 wt.% NaCl solution. A consistent finding across all specimens is that pit nucleation was exclusively associated with non-metallic inclusions at the matrix/inclusion interface. In both as-built and sensitized SLMed samples, pits preferentially nucleate at the matrix/inclusion interfaces of sulfide inclusions (Figures 7(a1) and (b1)) and composite sulfide-oxide inclusions (Figures 7(a2) and (b2)). Similarly, metastable pitting in as-received and sensitized wrought samples also originates from the same types of interfaces, specifically at sulfide inclusions (Figures 8(a1) and (b1)) and composite sulfide-oxide inclusions (Figures 8(a2) and (b2)). Notably, no pit initiation is observed at pure oxide inclusions.

Figure 7.

Typical SEM micrographs and corresponding EDS spectra of pit initiation sites in (a1, a2) as-built and (b1, b2) sensitized SLMed 316L SSs: (a1, b1) sulfide inclusions, (a2, b2) composite sulfide-oxide inclusions.

Figure 8.

Typical SEM micrographs and corresponding EDS spectra of pit initiation sites in (a1, a2) as-received and (b1, b2) sensitized wrought 316L SSs: (a1, b1) sulfide inclusions, (a2, b2) composite sulfide-oxide inclusions.

Figure 9 presents the Nyquist and Bode plots of the two 316L SSs measured at OCP in 3.5 wt.% NaCl solution. As shown in Figures 9(a1) and (b1), the Nyquist plots display depressed semicircles characterized by capacitive arcs. Similar capacitive behavior is observed across all samples throughout nearly the entire frequency range, suggesting that sensitization does not alter the fundamental corrosion mechanism in either SLMed or wrought specimens. The diameter of the capacitive semi-circle is associated with the corrosion resistance of the passive films.³⁶ As shown in Figures 9(a1) and (b1), the capacitive semi-circle diameters for both SLMed and wrought 316L SSs decrease after sensitization, demonstrating reduced corrosion resistance of their passive films. Notably, the SLMed 316L exhibits a larger semi-circle diameter than its wrought counterpart. Impedance analysis reveals a single time constant (Figures 9(a2) and (b2)), where log|Z| varies linearly with log f over a broad frequency range. The as-built SLMed 316L shows phase angles θ approaching −80°, indicative of a highly stable passive film formed in 3.5 wt% NaCl solution.

Figure 9.

(a1, b1) Nyquist and (a2, b2) bode plots of passive films formed on (a1, a2) SLMed and (b1, b2) wrought 316L SSs in 3.5 wt.% NaCl solution

The equivalent circuit model used for fitting the EIS results is shown in Figure 9(a1). In this model, R_s represents the electrolyte resistance, R₁ denotes the charge transfer resistance, and Q₁ corresponds to the constant phase element (CPE) of the electric double layer. The fitted parameters are listed in Table 4. According to previous studies,³⁷ when the impedance response of the CPE correlates with the effective capacitance, the dielectric constant or film thickness can be determined using the following equation:

C_{e f f} = \frac{ε ε_{0}}{δ}

(3)

where δ is the film thickness,

ε

is the dielectric constant, and

ε_{0}

is the permittivity of vacuum. The film thickness can be obtained from the power-law model.³⁷ In this model, the effective capacitance is expressed as:

C_{e f f, P L} = g Q (ρ_{δ} ε ε_{0})^{1 - n}

(4)

where

ρ_{δ}

is the resistivity, assigned a value of 500 Ω·cm in this model,

ε

is the relative dielectric constant of the passive film, and g is a function of n, which can be expressed by the following formula:

g = 1 + 2.88 (1 - n)^{2.375}

(5)

Table 4.

The fitted electrochemical parameters from EIS analysis and calculated passive film thickness for SLMed and wrought 316L SSs.

316L SSs		R_s (Ω·cm²)	Q₁		R₁×10⁵ (Ω·cm²)	δ (nm)
316L SSs		R_s (Ω·cm²)	Y₀×10⁻⁵ (Ω⁻¹·cm⁻²·sⁿ)	n	R₁×10⁵ (Ω·cm²)	δ (nm)
SLMed	As-built	8.16	1.28	0.93	42.10	3.68
SLMed	Sensitized	8.36	2.54	0.91	7.37	2.83
Wrought	As-received	8.83	1.43	0.93	31.50	3.29
Wrought	Sensitized	8.56	6.13	0.92	2.97	0.95

Therefore, the passive film thickness of the experimental steels was calculated, with the results summarized in Table 4. The electrolyte resistance shows minimal variation across the experimental steels, indicating no significant change in electrolyte composition during film formation. After sensitization, the charge transfer resistance decreases from 4.21 × 10⁶Ω cm² to 7.37 × 10⁵ Ω cm² for SLMed 316L SS, and from 3.15 × 10⁶ Ω cm² to 2.97 × 10⁵ Ω cm² for wrought 316L SS, suggesting diminished protective capability of the passive film. Notably, the wrought sample consistently exhibits lower protective capability than the SLMed sample. Furthermore, the passive film thickness was reduced by sensitization for both 316L SSs, but the SLMed 316L SS still retained a thicker film than the wrought sample after sensitization.

Figure 10 shows the Mott-Schottky plots of passive films formed on SLMed and wrought 316L SSs in 3.5 wt.% NaCl solution. Two linear regions with positive and negative slopes are observed in the Mott-Schottky plots, indicating that the passive films possess a dual-layer structure consisting of an outer layer rich in iron oxides and an inner layer enriched with chromium oxides.³⁸ In the positive-slope region I, the passive film exhibits n-type semiconductor characteristics, indicating that oxygen vacancies and/or cation interstitials are the major defect in the film.³⁹ Region II, characterized by a negative slope, shows p-type semiconductor behavior, with cation vacancies acting as the predominant defects.

Figure 10.

The Mott-Schottky plots of the passive films formed on (a) SLMed and (b) wrought 316L SSs in 3.5 wt.% NaCl solution.

The calculated values of N_D and N_A based on Eqs. (1) and (2) are summarized in Table 5. These parameters reflect the conductivity of the passive film, as higher N_D and N_A values correspond to greater conductivity, which increases the probability of film breakdown and pitting initiation. The magnitudes of N_A and N_D fall within the range of 10²⁰∼10²² cm⁻³, consistent with the values previously reported for passive films on stainless steels. These results reveal the highly disordered nature of the passive film structure.^39–41 It is evident that the acceptor densities are consistently higher than the donor densities. After sensitization, both SLMed and wrought 316L SSs show increased doping densities. Furthermore, the wrought 316L SS maintains higher doping densities in its passive film compared to the SLMed counterpart. Consequently, the deterioration in corrosion resistance for both materials after sensitization is attributable to alterations in passive film properties.

Table 5.

Electronic properties (N_D-donor density, N_A-acceptor density) of passive films formed on SLMed and wrought 316L SSs.

316L SSs	SLMed		Wrought
316L SSs	As-built	Sensitized	As-received	Sensitized
N_D (10²¹ cm⁻³)	0.70	1.93	1.52	5.73
N_A (10²¹ cm⁻³)	1.86	2.83	2.55	6.49

Figure 11 shows the XPS spectra of SLMed and wrought 316 L SSs before and after sensitization. As shown in Figures 11(a1) and (b1), the Cr2p_3/2 spectra from the passive films were deconvoluted into three constituent peaks corresponding to metallic Cr⁰ (574.0 eV), Cr₂O₃ (576.6 eV) and Cr(OH)₃ (577.7 eV). In passive films formed on stainless steels, Cr(OH)₃ and Cr₂O₃ are known to predominantly occupy the outer and inner layers, respectively.^42,43 The Fe2p_3/2 spectra presented in Figures 11(a2) and (b2) exhibit characteristic peaks for metallic Fe (706.1 eV), FeO (707.8 eV), Fe₂O₃ (710.4 eV) and FeOOH (711.8 eV). All measured binding energies show good agreement with reference values reported in the literature.¹⁵

Figure 11.

XPS spectra of the Cr2p_3/2 and Fe2p_3/2 detected for passive films formed on (a1, a2) SLMed and (b1, b2) wrought 316L SSs: (a1, b1) Cr2p_3/2, (a2, b2) Fe2p_3/2.

Based on the XPS spectra, the relative fraction of atomic i and j in passive film can be calculated using the following equation⁴⁴:

\frac{n_{i}}{n_{j}} = \frac{I_{i}}{I_{j}} \times \frac{σ_{j}}{σ_{i}} \times \frac{E_{k_{j}}^{0.5}}{E_{k_{i}}^{0.5}}

(6)

where n is the amount of surface atomics, I is the integrated peak area of the corresponding spectra, σ is the photoionisation cross-section of the relevant element's energy level, and E_k is the kinetics energy of photoelectron. In the present study, the σ values for Cr2p_3/2 and Fe2p_3/2 were taken as 7.69 and 10.82, respectively. The passive films are primarily composed of oxides and hydroxides of Fe and Cr, including FeO, Fe₂O₃, FeOOH, Cr₂O₃ and Cr(OH)₃. The relative concentrations of Cr and Fe were calculated from the Cr2p_3/2 and Fe2p_3/2 spectra. The ratios of {[Fe_ox]+ [Fe_hyd]}/{[Cr_ox] + [Cr_hyd]} in the passive films of as-built SLMed and wrought 316L SSs are 0.94 and 0.61, respectively. After sensitization, these ratios increase to 1.43 (SLMed) and 1.81 (wrought), indicating higher iron content in the passive film of the SLMed sample compared to the wrought counterpart, with both showing elevated iron levels after sensitization. The ratios of {[Fe_ox] + [Cr_ox]}/{[Fe_hyd] + [Cr_hyd]} are 1.25 (SLMed) and 2.51 (wrought), which slightly decrease to 1.21 and 2.19 after sensitization. The results demonstrate that the passive film of SLMed 316L SS contains a higher proportion of hydroxides than that of the wrought sample, and the Cr-oxide content decreases after sensitization in both materials. According to the bilayer theory of passive films on stainless steels, Cr₂O₃ is the primary constituent of the inner layer, while iron oxides and hydroxide, along with Cr(OH)₃ predominantly form the outer layer.⁴⁵ The XPS results indicate that the iron content in the outer passive film layer is significantly higher in SLMed 316L SS compared to its wrought counterpart, and the Cr-oxide content in the inner layer decrease after sensitization for both materials.

Discussion

This study systematically investigates the corrosion behavior of SLMed 316L SS after sensitization. The SLMed material exhibits significantly higher susceptibility to IGC, yet demonstrates superior pitting resistance compared to its wrought counterpart. This apparent contradiction is attributed to their fundamental microstructural differences. Through a comprehensive analysis combining EBSD, electrochemical measurements, and surface characterization, this work provides a unified mechanistic interpretation from three perspectives: GB character, defect-mediated passivation, and inclusion refinement.

The inferior IGC resistance of the SLMed 316L SS is attributed to its unfavorable GB character distribution. EBSD analysis confirms that the proportion of Σ3 CSL boundaries in the SLMed sample (9.71%) is significantly lower than that in the wrought counterpart (18.90%). During sensitization, Cr-rich carbides preferentially precipitate at HAGBs due to their inherently higher interfacial energy and enhanced diffusion rates, resulting in the formation of localized Cr-depleted zones. In contrast, the coherent Σ3 CSL boundaries effectively resist precipitation. As low-energy boundaries, they not only inhibit the formation of Cr-rich phases but also disrupt the interconnectivity of the HAGB network.^46–48 Consequently, the reduced density of Σ3 CSL boundaries in the SLMed 316L SS facilitates extensive precipitation along interconnected HAGBs, which compromises its IGC resistance. This mechanism is conclusively demonstrated by DL-EPR measurements and the corresponding intergranular attack morphologies.

The superior pitting resistance of the SLMed 316L SS is attributed to two synergistic mechanisms, with the substantial refinement of inclusions serving as the primary factor. Statistical analysis confirms that the rapid solidification inherent to the SLM process effectively reduces the population and size of inclusions. Crucially, SEM/EDS analysis demonstrates that all observed metastable pits nucleate exclusively at sulfide-containing inclusions. This nucleation selectivity is governed by the intrinsic dissolution kinetics of different inclusion types: sulfide inclusions undergo preferential dissolution in chloride environments, generating aggressive micro-crevices that concentrate chloride ions, whereas oxide inclusions (e.g., Al₂O₃) remain electrochemically inert and exhibit no catalytic activity for pit initiation.^49,50 Thus, sulfide-containing inclusions act as the primary initiation sites. Consequently, the refinement of these active sulfide-containing inclusions fundamentally eliminates the primary sites for pitting initiation.

A pivotal finding of this work is the beneficial role of the high-density defect structures — specifically, the elevated GB density and dislocation networks — in the SLMed sample, which synergistically enhance pitting resistance by promoting a superior passive film. This finding challenges the classical view that predominantly considers such defects as potential corrosion initiation sites.⁵¹ Instead, we demonstrate that these high-energy sites, including both GBs and dislocations, serve as preferential nucleation sites for passive film formation and as short-circuit diffusion pathways for Cr and Mo, significantly accelerating passivation kinetics.^52,53 This defect-mediated formation mechanism results in a passive film with altered composition and structure. Although XPS analysis indicates a higher Fe/Cr ratio in the passive film of the SLMed sample (0.94) compared to the wrought counterpart (0.61), the film shows a markedly higher hydroxide content. While literature suggests that a high Cr/Fe ratio is favorable for optimal film stability,^54–56 our results reveal that in the SLMed 316L SS, the detrimental impact of a moderately higher Fe/Cr ratio is effectively compensated by the rapid formation of a thick, hydroxide-rich film enabled by high-density defects. This unique microstructure thus promotes the development of a passive layer with robust barrier properties, as electrochemically evidenced by its higher charge-transfer resistance (EIS results) and lower donor density (Mott-Schottky analysis).

Sensitization treatment significantly deteriorates the pitting corrosion resistance of both materials, albeit to different degrees. This degradation mechanism originates from the precipitation of Cr-rich phases, which depletes the matrix of chromium and consequently leads to inferior passive film quality, as evidenced by XPS analysis showing increased Fe/Cr ratios (1.43 for SLMed, 1.81 for wrought) and reduced Cr-oxide content. This, in turn, facilitates pitting initiation (shown by the increased nucleation frequency in Figure 6) and accelerates the metal dissolution rate within active pits (evidenced by severe cumulative damage in sensitized samples). However, due to its inherent advantages of refined inclusions and an excellent repassivation capability enabled by high-density defect structures, the SLMed sample maintains relatively superior pitting resistance even after sensitization. When metastable pits initiate at dissolved inclusions, the highly protective passive film of the surrounding matrix, enhanced by these defect structures, effectively suppresses their transition into stable pits.⁵⁷ This accounts for the lower cumulative damage observed for the SLMed sample during potentiostatic tests.

It should be noted that this study employed a standardized 3.5 wt.% NaCl solution (room temperature, neutral pH) as a benchmark environment. This approach was essential to isolate and elucidate the fundamental influence of microstructural characteristics on corrosion behavior by eliminating interference from complex environmental variables. The identified pitting mechanisms of inclusion-assisted nucleation and passive film quality are universally recognized as fundamental to the pitting behavior of stainless steels in chloride-containing media.^58,59 Therefore, the superior performance of SLMed 316L SS in this benchmark system strongly indicates an inherent robustness that is likely to be maintained in more aggressive environments, such as those with elevated temperatures or altered pH. Future work will systematically investigate the coupled effects of these environmental parameters to evaluate the material's suitability for specific industrial applications.

Conclusions

The intergranular corrosion (IGC) susceptibility, electrochemical characteristics, and passive behavior of sensitized selective laser melted (SLMed) and wrought 316L SSs were systematically investigated using electron backscatter diffraction (EBSD), electrochemical measurements, and surface analysis. The main conclusions can be drawn as follows:

Double-loop electrochemical potentiokinetic reaction (DL-EPR) tests and corresponding corrosion microscopy reveal that SLMed 316L SS exhibits higher susceptibility to IGC than its wrought counterpart. This behavior is primarily attributed to its distinct grain boundary character distribution, particularly the significantly lower proportion of ∑3 coincident-site lattice (CSL) boundaries.

Electrochemical measurements demonstrate that sensitization compromises the pitting corrosion resistance of both SLMed and wrought 316L SSs. Nevertheless, the SLMed 316L SS exhibits a higher pitting potential than the wrought counterpart, and this is ascribed to the refined inclusion distribution in the SLMed material.

For both SLMed and wrought 316L SSs in both non-sensitized and sensitized conditions, metastable pits consistently nucleate preferentially at inclusion/matrix interfaces. Sensitization significantly compromises the protective capacity of passive films on both materials, resulting in increased cumulative charge values and higher initiation frequency of metastable pitting events in sensitized samples.

Sensitization reduces both the chromium oxide content and passive film thickness on both materials, thereby reducing their protective capability. Compared to the wrought material, the SLMed 316L SS exhibits higher hydroxide content and thicker passive films, which can be attributed to its higher grain boundary density and dislocation network.

Footnotes

ORCID iD

Yan Hou

Author contribution(s)

Yan Hou: Conceptualization; Data curation; Formal analysis; Investigation; Writing – original draft.

Jin-Zhe Xia: Data curation; Formal analysis; Investigation; Writing – original draft.

Ke-Qing Zhang: Data curation; Formal analysis; Investigation.

Nan Zhang: Data curation; Formal analysis; Investigation.

Chen-Yu Li: Formal analysis; Project administration; Resources; Writing – review & editing.

Chao-Long Li: Formal analysis; Project administration; Resources; Writing – review & editing.

Xiao-Yun Yuan: Methodology; Project administration; Resources.

Chun-Fang Ma: Project administration; Supervision.

Funding

The authors disclosed receipt of the following financial support for the research, authorship, and/or publication of this article: This work is supported by Natural Science Basic Research Program of Shaanxi Province (2022JQ-453, 2022JQ-421, 2023-JC-QN-0462); National College Students Innovation and Entrepreneurship Training Program (S202310709054); PHD Start-up Fund of XPU (107020620), National Sustainable Development Agenda Innovation Demonstration Zone Construction Science and Technology Project of ChengDe, Hebei(202302F011).

Declaration of conflicting interests

The authors declared no potential conflicts of interest with respect to the research, authorship, and/or publication of this article.

Data availability statement

The data that support the findings of this study are available on request from the corresponding author.

References

Kale

Kim

, et al. An investigation of the corrosion behavior of 316L stainless steel fabricated by SLM and SPS techniques. Mater Charact 2020; 163: 110204.

Wang

, et al. Control of shape and performance for direct laser fabrication of precision large-scale metal parts with 316L stainless steel. Opt Laser Technol 2013; 45: 209–216.

Sander

Thomas

Cruz

, et al. On the corrosion and metastable pitting characteristics of 316L stainless steel produced by selective Laser melting. J Electrochem Soc 2017; 164: 250–257.

Fujii

Suzuki

Shimamura

. Susceptibility to intergranular corrosion in sensitized austenitic stainless steel characterized via crystallographic characteristics of grain boundaries. Corros Sci 2022; 195: 109946.

Tokuda

Muto

Sugawara

, et al. Effect of sensitization on pitting corrosion at MnS and CrS in Type 304 stainless steel. J Electrochem Soc 2021; 168: 091504.

Cheng

Klinkenberg

Ise

, et al. Pitting corrosion of sensitised type 304 stainless steel under wet-dry cycling condition. Corros Sci 2017; 118: 217–226.

Voisin

Shi

Zhu

, et al. Pitting corrosion in 316L stainless steel fabricated by Laser powder bed fusion additive manufacturing: a review and perspective. J Miner Met Mater Soc 2022; 74: 1668–1689.

Yap

Chua

Dong

, et al. Review of selective laser melting: materials and applications. Appl Phy Rev 2015; 2: 1–22.

Zou

Liu

Yang

, et al. Corrosion behavior of different building planes of selective laser melting 316L stainless steel in 0.1 M HCl solution. J Mater Res Technol 2024; 28: 4738–4753.

10.

Liverani

Toschi

Ceschini

, et al. Effect of selective Laser melting (SLM) process parameters on microstructure and mechanical properties of 316L austenitic stainless steel. J Mater Process Technol 2017; 249: 255–263.

11.

Pokharel

Balogh

Brown

, et al. Signatures of the unique microstructure of additively manufactured steel observed via diffraction. Scr Mater 2018; 155: 16–20.

12.

Pham

Dovgyy

Hooper

. Twinning induced plasticity in austenitic stainless steel 316L made by additive manufacturing. Mater Sci Eng A 2017; 704: 102–111.

13.

Fang

Huang

, et al. Review on residual stress in selective laser melting additive manufacturing of alloy parts. Opt Laser Technol 2020; 129: 106283.

14.

Wang

Zou

Qiao

, et al. Influence of selective laser melting process parameters on the microstructure, corrosion resistance, and mechanical properties of 2205 duplex stainless steel. J Mater Res Technol 2025; 38: 1825–1841.

15.

Man

Dong

Liu

, et al. The enhancement of microstructure on the passive and pitting behaviors of selective laser melting 316L SS in simulated body fluid. Appl Surf Sci 2019; 467–468: 193–205.

16.

Chen

Cheng

, et al. Effect of heat treatment on microstructure, mechanical and corrosion properties of austenitic stainless steel 316L using arc additive manufacturing. Mater Sci Eng A 2018; 715: 307–314.

17.

Man

Duan

Cui

, et al. The effect of sub-grain structure on intergranular corrosion of 316L stainless steel fabricated via selective laser melting. Mater Lett 2019; 243: 157–160.

18.

Macatangay

Thomas

Birbilis

, et al. Unexpected interface corrosion and sensitization susceptibility in additively manufactured austenitic stainless steel. Corrosion 2017; 74: 153–157.

19.

Laleh

Hughes

, et al. On the unusual intergranular corrosion resistance of 316L stainless steel additively manufactured by selective laser melting. Corros Sci 2019; 161: 108189.

20.

Snitzer

Lou

. Sensitization of 316L stainless steel made by Laser powder bed fusion additive manufacturing. Corrosion 2023; 79: 240–251.

21.

Macatangay

DAT

Conrades

Brunner

, et al. Applications of wrought austenitic stainless steel corrosion testing to Laser powder bed fusion 316L. Corrosion 2021; 78: 13–24.

22.

Taji

Moayed

Mirjalili

. Correlation between sensitisation and pitting corrosion of AISI 403 martensitic stainless steel. Corros Sci 2015; 92: 301–308.

23.

Shit

Mariappan

Ningshen

. Improvement of sensitization and intergranular corrosion of AISI type 304L stainless steel through thermo-mechanical treatment. Corros Sci 2023; 213: 110975.

24.

Liang

Dong

, et al. Influence of sensitization on microstructure and passive property of AISI 2205 duplex stainless steel. Corros Sci 2016; 104: 144–151.

25.

Ida

Muto

Sugawara

, et al. Local electrochemistry and in situ microscopy of pitting at sensitized grain boundary of Type 304 stainless steel in NaCl solution. J Electrochem Soc 2017; 164: 779–787.

26.

Özer

. Heat treatment, microstructure and corrosion relationship on 316L SS produced by AM. Mater Sci Technol 2023; 39: 671–682.

27.

Zhang

Zhao

, et al. Effect of heat treatments on metastable pitting of 316L stainless steel fabricated by selective laser melting. J Mater Res Technol 2022; 21: 1903–1914.

28.

Kain

Chandra

Adhe

, et al. Effect of cold work on low-temperature sensitization behaviour of austenitic stainless steels. J Nucl Mater 2004; 334: 115–132.

29.

Guo

Liang

. The effect of pre-deformation on corrosion resistance of the passive film formed on 2205 duplex stainless steel. J Alloys Compd 2016; 686: 176–183.

30.

Hsu

Szpunar

, et al. Deformation mechanism and texture and microstructure evolution during high-speed rolling of AZ31B Mg sheets. J Mater Sci 2008; 43: 7148–7156.

31.

Saeidi

Ashrafizadeh

Niroumand

, et al. EBSD Study of micromechanisms involved in high deformation ability of DP steels. Mater Des 2015; 87: 130–137.

32.

Heine

Kirchheim

. Dissolution rates of iron and chromium and Fe-Cr alloys in the passive state. Corros Sci 1990; 31: 533–538.

33.

Soltis

. Passivity breakdown, pit initiation and propagation of pits in metallic materials. Rev Corr Sci 2015; 90: 5–22.

34.

Dong

Guo

Zheng

, et al. Investigation on inhibition of CrO₄²⁻and MoO₄²⁻ ions on carbon steel pitting corrosion by electrochemical noise analysis. J Appl Electrochem 2002; 32: 395–400.

35.

Bai

, et al. Effects of boron on microstructure and metastable pitting corrosion behavior of Super304H austenitic stainless steel. J Electrochem Soc 2015; 162: 473–481.

36.

Cao

. On the impedance plane displays for irreversible electrode reactions based on the stability conditions of the steady-state—I. One state variable besides electrode potential. Electrochim Acta 1990; 35: 831–836.

37.

Orazem

Frateur

Tribollet

, et al. Dielectric properties of materials showing constant-phase-element (CPE) impedance response. J Electrochem Soc 2013; 160: 215–225.

38.

Gupta

Birbilis

. The influence of nanocrystalline structure and processing route on corrosion of stainless steel: a review. Corros Sci 2015; 92: 1–15.

39.

Taveira

Montemor

Belo

MDC

, et al. Influence of incorporated Mo and Nb on the Mott-Schottky behaviour of anodic films formed on AISI 304L. Corros Sci 2010; 52: 2813–2818.

40.

Hakiki

Montemor

Ferreira

MGS

, et al. Semiconducting properties of thermally grown oxide films on AISI 304 stainless steel. Corros Sci 2000; 42: 687–702.

41.

Luo

Dong

, et al. The electrochemical behaviour of 2205 duplex stainless steel in alkaline solutions with different pH in the presence of chloride. Electrochim Acta 2012; 64: 211–220.

42.

Guo

Liu

, et al. Effect of passive film on mechanical properties of martensitic stainless steel 15-5PH in a neutral NaCl solution. Appl Surf Sci 2015; 327: 313–320.

43.

Cui

Chen

Wang

, et al. Passivation behavior and surface chemistry of 2507 super Duplex stainless steel in acidified artificial seawater containing thiosulfate. J Electrochem Soc 2017; 164: 856–868.

44.

Cheng

Zhao

Cao

, et al. Facile chromaticity approach for the inspection of passive films on austenitic stainless steel. Corros Sci 2013; 70: 235–242.. http://

45.

Man

Dong

Cui

, et al. A comparative study of primary and secondary passive films formed on AM355 stainless steel in 0.1 M NaOH. Appl Surf Sci 2018; 427: 763–773.

46.

Yin

Song

Deng

, et al. Thermal stability and microstructural evolution of additively manufactured 316L stainless steel by laser powder bed fusion at 500-800°C. Addit Manuf 2021; 41: 101981.

47.

Yang

Liu

Song

, et al. Sensitization, desensitization, and carbide evolution of Alloy 800H made by laser powder bed fusion. Additive Manufacturing 2022; 50: 102547.

48.

Padilha

Escriba

Materna-Morris

, et al. Precipitation in AISI 316L(N) during creep tests at 550 and 600°C up to 10 years. J Nucl Mater 2007; 362: 132–138.

49.

Sun

Tan

Lan

, et al. Mechanisms of inclusion-induced pitting of stainless steels: a review. J Mater Sci Technol 2024; 168: 143–156.

50.

Takayama

Nishimoto

Muto

, et al. Micro-Electrochemical aspects of the effects of temperature on pit initiation at MnS inclusion in type 304 stainless steel. J Electrochem Soc 2022; 169: 111501.

51.

Örnek∗

Engelberg

. SKPFM Measured volta potential correlated with strain localisation in microstructure to understand corrosion susceptibility of cold-rolled grade 2205 duplex stainless steel. Corros Sci 2015; 99: 164–171.

52.

Gaberšček

Pejovnik

. Impedance spectroscopy as a technique for studying the spontaneous passivation of metals in electrolytes. Electrochim Acta 1996; 41: 1137–1142.

53.

Sun

Xiao

Dong

, et al. Effect of stress on electrochemical characteristics of pre-cracked ultrahigh strength stainless steel in acid sodium sulphate solution. Corros Sci 2014; 89: 137–145.

54.

Yue

Chen

Krishnan

, et al. In depth analysis of the passive film on martensitic tool alloy: effect of tempering temperature. Corros Sci 2024; 234: 112133.

55.

Liu

Gong

Yang

, et al. Electrochemical dissolution behavior of stainless steels with different metallographic phases and its effects on micro electrochemical machining performance. Electrochem Commun 2024; 160: 107677.

56.

Abreu

Cristóbal

Losada

, et al. Comparative study of passive films of different stainless steels developed on alkaline medium. Electrochim Acta 2004; 49: 3049–3056.

57.

Ilevbare

Burstein

. The role of alloyed molybdenum in the inhibition of pitting corrosion in stainless steels. Corros Sci 2001; 43: 485–513.

58.

Hou

Nakamori

Kadoi

, et al. Initiation mechanism of pitting corrosion in weld heat affected zone of duplex stainless steel. Corros Sci 2022; 201: 110278.

59.

Che

Revilla

, et al. Role of te-RE alloying on the passive film and pitting corrosion behavior of 316L stainless steel. Corros Sci 2024; 240: 112457.