Pitting behavior of SLM 316L stainless steel exposed to chloride environments with different aggressiveness: Pitting mechanism induced by gas pores

Journal Pre-proof Pitting behavior of SLM 316L stainless steel exposed to chloride environments with different aggressiveness: pitting mechanism induced by gas pores Zhiwei Duan (Conceptualization) (Methodology) (Data curation), Cheng Man (Validation) (Supervision) (Writing - original draft), Chaofang Dong (Writing - review and editing), Zhongyu Cui (Data curation) (Writing - review and editing), Decheng Kong (Software) (Investigation), Li wang (Visualization) (Data curation), Xin Wang (Project administration)

PII:

S0010-938X(19)31936-5

DOI:

https://doi.org/10.1016/j.corsci.2020.108520

Reference:

CS 108520

To appear in:

Corrosion Science

Received Date:

15 September 2019

Revised Date:

28 January 2020

Accepted Date:

4 February 2020

Please cite this article as: Duan Z, Man C, Dong C, Cui Z, Kong D, wang L, Wang X, Pitting behavior of SLM 316L stainless steel exposed to chloride environments with different aggressiveness: pitting mechanism induced by gas pores, Corrosion Science (2020), doi: https://doi.org/10.1016/j.corsci.2020.108520

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Pitting behavior of SLM 316L stainless steel exposed to chloride environments with different aggressiveness: pitting mechanism induced by gas pores Zhiwei Duan a, Cheng Man a, b , Chaofang Dong b, Zhongyu Cui a, Decheng Kong b, Li wang b, Xin Wang a a.

School of Materials Science and Engineering, Ocean University of China, Qingdao 266100 China Beijing Advanced Innovation Center for Materials Genome Engineering, Key

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b.

Laboratory for Corrosion and Protection (MOE), Institute for Advanced Materials and Technology, University of Science and Technology Beijing, Beijing 100083,

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China

Corresponding Author

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E-mail: [email protected]

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Highlights

Pitting corrosion of SLM 316L SS in different aggressive solutions was studied.

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A model was built to explain the pitting mechanism induced by gas-pore.

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Metastable pitting at covered gas-pores is more likely to transition into stable.

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Abstract

In this study, the pitting behaviors of wrought and SLM 316L SSs (stainless steels)

are comparatively investigated under the conditions of different aggressiveness. The experimental results show that SLM 316L SS exhibits higher sensitivity to pitting corrosion in extremely aggressive solutions, while the wrought sample is more

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vulnerable to pitting in conditions with low aggressiveness. Gas pores are the pittingsensitive sites in SLM 316L SS, and metastable pits initiated at covered gas pores are easier to evolve to stable growth compared to open pores due to the greater diffusion resistance. Keywords: A. Stainless steel, B. EIS, B. Mott-Schottky, C. pitting corrosion, C. Passive film

1. Introduction

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Selective laser melting (SLM), which is a new additive manufacturing technology, can efficiently fabricate metallic components with complex shapes [1, 2]. This

technology has been well established for 316L austenitic stainless steel [3-5]. In general,

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the superior corrosion resistance of 316L stainless steel is a crucial factor for its extensive application. Hence, for the popularization of SLM 316L SS (316L stainless

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steel fabricated by SLM), it is necessary to investigate its corrosion behavior. The manufacturing process of SLM technique is completely different from that of

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conventional forging, which results in significant differences in microstructures between the SLM and wrought 316L SSs [6-8]. First, pores, which are formed by entrapped gas or lack of fusion of powder particles in the SLM build, are one of the

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most typical microstructures in SLM 316L SS [9-11]. Second, solute enrichment of the liquid in front of the solid/liquid interface of the SLM melting pools, which determines

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the constitutional undercooling, can induce dendritic growth during the solidification of SLM 316L SS, as reported by Casalino et al. [12]. From previous literature, dendritic

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growth not only gives rise to the formation of fine subgrain in an individual grain [13, 14] but also creates the irregular shapes of the grains in the SLM 316L SS [10]. Due to the fine subgrains and large irregular grains, the interface structures in the SLM 316L SS are more complex compared with the wrought 316L SS [4]. Third, the estimated cooling rate during the SLM process can reach up to 105 K/s – 107 K/s, a value much higher than ingot casting (approximately 1 – 100 K/s), which results in the dense dislocations and high residual stresses in the SLM 316L SS [8, 14, 15]. Kong et al. 2

reported that the dislocation and residual stress were enriched at the grain and subgrain boundaries [3]. Finally, the size of the inclusions in SLM 316L SS was much smaller than in the wrought sample, due to the limitation of the rapid cooling in the growth of inclusions. As reported by Chao et al., the spherical inclusions in SLM 316L SS had diameters less than 100 nm and were composed of the elements Cr, Si, Mn and O [16]. Overall, pores, subgrains, irregular grains, dense dislocations and fine inclusions are the main typical microstructures of SLM 316L SS distinguishing it from wrought 316L SS. It is well known that the corrosion behavior of metallic materials is closely related

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to their microstructures. Therefore, unraveling the influence of these typical microstructures on the passivation and pitting corrosion is beneficial for understanding the corrosion behavior of SLM 316L SS.

Recently, to illuminate the relationship among typical microstructures of SLM

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stainless steel and their corrosion behavior, the passivation and pitting corrosion of SLM stainless steels were compared with these of stainless steels fabricated by

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conventional methods. Schindelholz et al. investigated the corrosion kinetics of SLM and wrought 304 SSs immersed in four acidic solutions and reported that SLM 304 SS

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possessed lower corrosion rates in 5% HCl + 1000 ppm FeCl3 and boiling 30% HNO3 solutions, but higher corrosion rates in 5% HCl and boiling 99% acetic acid solutions,

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compared to the wrought specimens [17]. A similar phenomenon was obtained in previous works of our group. When estimating the corrosion resistance of the wrought and SLM 316L SSs in simulated body fluid (SBF) and simulated proton exchange

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membrane fuel cell (PEMFC) solution using electrochemical methods, we found that the SLM 316 SS was superior to the wrought sample in terms of electrochemical

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resistance in SBF, while the SLM specimen exhibited a larger current density than the wrought material with polarization at 0.6 VSCE in the other solution [3, 10, 18]. Based on these experimental results, it could be deduced that the influence of the typical microstructures on the corrosion behavior of SLM stainless steels should rely on the external environment, but this hypothesis must be further confirmed. In a study examining the pitting corrosion of SLM 316L SS in solution with 0.6 M NaCl, Chao et al. found that all the SLM 316L SS exhibited a higher pitting potential than the wrought 3

material, which was attributed to the absence of MnS inclusions in the SLM 316L SS that usually act as pitting sensitive sites in the wrought sample [16]. Sander et al. investigated the pitting corrosion of SLM 316L SSs with different porosities and reported that metastable pitting corrosion on SLM samples increased with their porosities, but the pitting potentials of all SLM 316L SSs were higher than the wrought specimen [5]. In a previous work, we confirmed that the pores were the pitting sensitive sites in SLM 316L SS by comparing the morphologies of some marked pores before and after potentiodynamic polarization, which was in good agreement with Sander’s

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results [10]. However, the process and mechanism of pitting corrosion induced by pores in SLM stainless steels were still unclear. According to Frankel and Scully, the process controlling pitting corrosion is related to the experimental conditions, that is, pit

stability considerations are controlled under aggressive conditions (harsh electrolytes

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and extreme environments and/or susceptible microstructures), and passive film properties and protectiveness are critical factors in less extreme environments and/or

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for less susceptible alloys [19]. In terms of this perspective, we attempted to explain the pitting mechanism initiated by pores via investigating the passivation and pit growth

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process of SLM 316L SS in environments with different harshness properties. In this paper, the corrosion behaviors of SLM and wrought 316L SSs were

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comparatively studied in solutions with different pH values and Cl- concentrations. The pitting potentials of the two kinds of materials were determined by potentiodynamic polarization, the properties of passive films were characterized using electrochemical

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impedance spectra (EIS) and Mott-Schottky and X-ray photoelectron spectroscopy (XPS), and the stable pit growth was investigated by an immersion test in 6% FeCl3.

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Finally, the mechanism of pitting corrosion induced by the pores in SLM 316L SS was discussed based on the experimental results.

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2. Experimental 2.1 Materials and solution The AISI316L SS powder with a size ranging 15 – 45 μm was used to fabricate the SLM sample. The used powder had the following chemical composition (wt.%): 17.5 Cr, 10.4 Ni, 2.7 Mo, 1.2 Mn, 0.4 Si, 0.02 C,  0.02 P,  0.01 S and Fe balanced. The SLM manufacture process was performed on an EOS M280 in continuous mode

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with parameters of 200 W laser power, 800 mm/s scan speed and 120 μm layer thickness. After fabrication, the compactness of the SLM sample was measured with drainage, and the sample with compactness of approximately 98.6% was selected for this work.

In addition, the wrought 316L SS used in this work had a composition (wt.%) of 17.8

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Cr, 11.2 Ni, 2.2 Mo, 0.9 Mn, 0.6 Si, 0.03 C,  0.01 P,  0.01 S and Fe balanced. The asprinted sample was regarded as the object of this study, and the microstructure and

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corrosion behavior of the plane perpendicular to the building direction was investigated. The NaCl contents of the solutions for electrochemical measurements were 0.01

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M, 0.1 M, 0.5 M, 1 M and 3 M, and the pH values were 1, 3, 5, 7, 9, 11 and 13. For the immersion test, the solution was prepared with 6% FeCl3, pH 3. All the solutions were

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made from analytical grade reagents (NaCl, FeCl3·6H2O, HCl) and deionized water. 2.2 Microstructure analysis

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The metallographic structures of the wrought and SLM 316L SSs were observed with a scanning electron microscope (SEM, FEI Quanta 250) after mechanical

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polishing and chemical etching in the diluted aqua regia (HCl: HNO3: H2O = 3:1:4) for 40 s. The electron backscattered diffraction (EBSD, EDAX) equipped with the SEM was utilized to characterize the grain orientations of the two types of stainless steels after electropolishing with 20% perchloric acid in liquid nitrogen at 20 V for approximately 30 s, and the EBSD data were analyzed using TSL OIM Analysis 7 software. The submicro microstructural features, including subgrain boundaries, dislocations and inclusions, were obtained with a transmission electron microscope 5

(TEM, JEOL 2100), and the composition of the inclusions was analyzed by energy disperse spectroscopy (EDS). The samples for TEM observation were polished with diamond paste and finally thinned by ion-milling to achieve electron transparency. 2.3 Electrochemical measurements The electrochemical measurements were carried out on an electrochemical workstation (Autolab PGSTAT 302N) in the three-electrode system: the wrought/SLM 316L SS was the working electrode, a platinum sheet with a size of 20 mm × 20 mm

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was the counter electrode, and a saturated calomel electrode (SCE, 0.242 V vs standard

hydrogen electrode) was the reference electrode. After being connected to a copper wire

on their backside, the samples used as the working electrode were embedded in epoxy resin, leaving a working area of 10 mm × 10 mm, and then wet ground sequentially to

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2000 grit SiC paper, degreased with alcohol, cleaned in water, and dried in cold air.

The cyclic potentiodynamic polarization scanned from -1.0 VSCE to a potential

and the point was associated with a sudden increase in current density, was regarded

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2

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with a scanning rate of 0.167 mV/s, where the anodic current density reached 1 mA·cm-

as the pitting potential (Epit). The potential was then reversed and progressed in the cathodic direction to a potential below the protection potential (Epro). Although the rapid

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increase was related to the oxygen evolution reaction in some conditions (such as 1 M NaCl, pH 13), the flexion point was still denoted by Epit for the sake of description.

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After the passive film formed by 1 h potentiostatic polarization at the potentials (-0.10, -0.05, 0, 0.05 and 0.10 VSCE) that were located at the passive region, the

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electrochemistry impedance spectrum (EIS) tests were performed in the frequency range from 100 kHz to 10 mHz, with 7 points per decade using an amplitude of 10 mV. The EIS data were analyzed using ZsimpWin software. Following the EIS tests, the Mott-Schottky measurements were conducted at 1000 Hz using an ac signal of 10 mV and a step rate of 50 mV/s. To assure repeatability, all the mentioned electrochemical measurements were performed at room temperature (approximately 25 ºC) and repeated at least 3 times. To induce pitting corrosion, another potentiostatic polarization at 0.25 6

VSCE that was close to Epit was carried out until the anodic current reached 0.6 mA·cm2

, and the time corresponding to a current density of 0.5 mA·cm-2 was considered the

pitting incubation time. The cyclic potentiodynamic polarization measurements were performed in solution (1 M NaCl, pH 1, 3, 5, 7, 9, 11 and 13, and pH 3 with 0.01 M, 0.1 M, 0.5 M, 1 M and 3 M), and the other electrochemical measurements, including potentiostatic polarization, EIS and Mott-Schottky, were carried out in 1 M NaCl solution, pH 3.

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2.4 Immersion test The immersion tests were carried out in a water bath kettle, where the temperature was controlled at 60 ºC. Prior to the test, the samples were embedded in epoxy resin

leaving a working area of 10 mm × 10 mm, and then polished and cleaned in accordance

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with the procedures described for the samples used in the electrochemical

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measurements. The period of the immersion test was 24 h, and three duplicate samples were included. In addition, another SLM sample with some marked gas pores was

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immersed for 48 h to research the different pitting sensitivity between different gas pores.

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2.4 Surface characterization

After formation via potentiostatic polarization at ±0.10 VSCE for 1 h, the

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compositions of the passive films were analyzed using X-ray photoelectron spectroscopy (XPS, Thermo ESCALAB 250Xi). All XPS peaks were fitted using

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Avantage software after correction to the standard carbon C1s with a binding energy of 285.0 eV.

The pitting morphologies formed during the immersion test were observed by

SEM. The pit sizes were measured by confocal laser scanning microscopy (CLSM, KEYENCE VK-X250), and then, the data obtained by CLSM were statistically analyzed.

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3. Results 3.1 Microstructure The metallographic structures of wrought and SLM 316L SSs are shown using the inverse pole figure (IPF) of EBSD. As shown in Fig. 1a and 1b, the wrought 316L SS possessed a typical recrystallized microstructure composed of equiaxed grains, while the SLM 316L had a structure similar to that of laser welding composed of columnar

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and cellular grains. Grains with wave-like boundaries were observed in SLM 316L SS, which differs from the wrought sample [13]. After mechanical polishing and chemical etching, an intragranular cellular segregation network structure was observed on the SLM 316L SS using SEM. The network structure was constituted by equiaxed and bar-

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like cells with an average width of approximately 500 nm, as shown in Fig. 1c and 1d.

This subgrain structure bore clear intragranular features because the adjacent units had

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almost the same crystallographic orientation, as displayed in Fig. 1b. According to the theory related to laser welding, the appearance of the cellular structures could be

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ascribed to the rapid cooling rate during SLM processes [20, 21]. The different shapes of subgrains indicated the different growth directions of the large grains containing them. The near equiaxed zone (zone ‘A’ marked in Fig. 1d) expressed the columnar-

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dendritic growth direction almost along the building direction (Z axis), and the elongated zone (zone ‘B’ marked in Fig. 1d) reflected a certain angle between the

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growth direction and the building direction. The TEM images (Fig. 1e and 1f) showed the submicro microstructural features

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of SLM 316L SS. Dislocations were concentrated at the subgrain boundaries, which confirmed the existence of thermal residual stress caused by the rapid cooling rate of the SLM process (Fig. 1e). Fig. 1f exhibits the spherical inclusions with sizes in the range of 50 – 200 nm, which were randomly distributed throughout the microstructure. The inclusion size in SLM 316L SS was several orders of magnitude smaller than the MnS inclusions with a size of 2.36 μm in the wrought 316L SS (shown in Fig. 1g and 1h). Assessment of the composition by EDS revealed the elements (Fe, Cr, Mn, Al, Si, 8

O and S) that were present in the inclusions. In terms of the quantitative analysis of EDS (as shown in the table inserted in Fig. 1f), the inclusions should be composed of metallic silicate, which is in good agreement with the results described by Birbilis [5, 16]. Shen et al. also reported that the inclusions in the SLM 316L SS were mainly composed of silicates, but they proposed that the content of Cr was much greater than other metallic elements in the silicate inclusions [8]. Despite some differences in the compositions of inclusions in the reported works, there were clearly no MnS inclusions detected in SLM 316L SS.

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Although the SLM technology of 316L SS has been remarkably enhanced over the past years, porosity is still a common defect in SLM-fabricated components. In the

SLM sample used in this experiment, two types of pores were detected using SEM: gas pores (Fig. 2) and lack of fusion pores (not presented here). The gas pores, which are

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characterized by circular striations, are due to entrapped gas during the production of gas atomized 316L SS powder, and the lack of fusion pores is caused by insufficient

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melting of metallic powder [22]. Furthermore, we found that the gas pores exhibited different morphologies after the samples were mechanically polished. If the cross-

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section was close to the bottom of the gas pores, flat and hemispherical defects without a cover were observed (Fig. 2a and 2b). Conversely, the emerged gas pores were

and 2d).

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covered by the remnant metal as the polished surface became closer to the top (Fig. 2c

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3.2 Pitting behavior of wrought and SLM 316L SSs

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The electrochemical processes and pitting corrosion resistance of wrought and SLM 316L SSs in the solutions with different aggressiveness were comparatively researched using potentiodynamic polarization curves. The aggressiveness of the solutions was adjusted by controlling the NaCl contents and pH values, where the more concentrated NaCl and lower the pH value, the higher was the aggressiveness for stainless steels [23, 24]. Fig. 3a and 3b plot the cyclic potentiodynamic polarization curves of the wrought and SLM 316L SSs in the solutions with different NaCl contents 9

(pH 3), respectively. No significant difference was detected in the shapes of these curves, suggesting that the two samples had similar electrochemical processes under these conditions. The anodic regions of the curves were characterized by a plateau of constant current densities followed by an abrupt current increase, indicating the passive and breakdown behaviors within the potential range [25]. Although the electrochemical processes of wrought and SLM 316L SSs reflected by the potentiodynamic polarization curves were similar, there are some differences in their pitting behavior. Taking the polarization curves of the two samples measured in 1 M NaCl solution at pH 3 as

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examples (Fig. 4a and 4b), the intensity and frequency of the anodic current fluctuations associated with metastable pitting of the wrought 316L SS were much larger than those of the SLM specimen. This result indicated that metastable pitting was more prominent in the wrought 316L SS compared to the SLM sample. In addition, the pitting potentials

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of these two types of stainless steels decreased with increasing NaCl content (Fig. 4c), but their variation tendencies with increasing pH value were different: the pitting

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potential of wrought 316L SS consistently became more positive, while the potential of the SLM sample increased first and then decreased (Fig. 4d). Based on the pitting

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potentials plotted in Fig. 4c and 4d, the pitting potential of the SLM 316L SS was slightly more positive than the wrought samples in all solutions except 1 M NaCl, pH

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1 and 3 M NaCl, pH 3, which could be regarded as the two most corrosive conditions. The different variations with the pH value of the pitting potentials of the wrought and SLM 316L SS should be due to their different pitting mechanisms. For the wrought

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316L SS, the pitting corrosion was always initiated at MnS inclusions, the stability of which increase with increasing pH; therefore, the pitting potential of wrought 316L SS

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increased with higher pH values. The initiation of pitting corrosion of the SLM 316L SS is related to the stability of the passive film formed on the bottom of the gas pores, and stainless steels always exhibit superior passive behavior in a neutral environment than in acidic and alkaline conditions [23]. The protection potentials of the SLM 316L SS were always more positive than those of the wrought sample in all the test solutions except that with a pH value of 13 (Fig. 4c and 4d), which indicated that repassivation was more likely to occur for the SLM 316L SS when the electrochemical active 10

dissolution was prevented. In addition, both the wrought and SLM 316L SSs had protection potentials that were nearly equal to their pitting potential in medium with a pH value of 13, which demonstrated that little pitting corrosion occurred and the rapid increase in current density in the polarization curves was related to the oxygen evolution reaction under this condition [26]. The variations of the pitting and protection potentials shown in Fig. 4c and 4d revealed that SLM 316L SS exhibited higher pitting resistance in the solution with slight aggressiveness, and less pitting resistance in the medium with strong aggressiveness, in contrast to the wrought 316L SS. This phenomenon confirmed

was dependent on the aggressiveness of the environment.

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that the influence of the typical microstructures on the corrosion behavior of SLM 316L

In previous papers, the difference in pitting behavior between wrought and SLM

316L SSs were related to the refinement/absence of MnS inclusions in the SLM 316L

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SS [5, 16]. However, the reduced pitting corrosion resistance of SLM 316L SS in the

medium with strong aggressiveness could not be explained using the absence of MnS

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or other inclusions. In a perspective paper, Frankel and Scully proposed that the pitting resistance of metals depends on the properties of passive film in less extreme

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environments and is related to the pit stability in aggressive conditions [19]. Therefore, the less pitting corrosion resistance of SLM 316L SS in strongly corrosive solutions

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will be discussed after the introductions regarding the properties of passive film (Section 3.3) and pit stability (Section 3.4).

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3.3 Passivation of wrought and SLM 316L SSs

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3.3.1 Passive film formation Fig. 5 displays the current vs time (i ~ t) curves in linear and double logarithmic

coordinates for wrought and SLM 316L SSs polarized at the selected anodic potentials for 1 h. From Fig. 5a and 5b, it can be seen that all the i ~ t curves had nearly identical shapes: the current density rapidly decreased with elapsing time in the initial stage of potentiostatic polarization, which is associated with the rapid nucleation and growth of the passive film [27]; then, the current density remained at a relative stable value, 11

indicating the balance of formation and dissolution of the passive film [28]. In addition, some current fluctuations could be clearly observed in the steady-current stages of the i ~ t curves of the wrought specimen polarized at 0.05 and 0.10 VSCE. As reported in previous papers, these current fluctuations are related to either nucleation or metastable pitting corrosion [29, 30]. As shown in Fig. 5c and 5d, linear regions are contained in the log-log coordinate. It has been proposed that the slopes of the linear regions can reflect the qualities of the passive films, where k = -1 denotes the formation of a compact film, while k = -0.5 indicates the presence of a porous passive film [28, 31].

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In this experiment, all the fitted slopes were close to -1, suggesting that the protective passive film could be formed in the selected film formation conditions. Finally, it is

noteworthy that the current density still declined at 1 h after polarization, proving the

formation of a quasi-steady-state, which was sufficient to attain the purpose of

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interpreting the different passive behavior of the wrought and SLM 316L SSs in this

3.3.2 Electrochemical impedance

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work [32, 33].

EIS measurements were performed to compare the relative stability of the passive

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films formed on wrought and SLM 316L SSs. Fig. 6 shows the EIS results depicted in the forms of Nyquist and Bode plots, where the points are the experimental data and

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the solid lines are the fitting results. The Nyquist diagrams of the EIS spectra display a similar shape that is characterized by a somewhat unfinished capacitance arc, demonstrating similar passive mechanism, as shown in Fig. 6a and 6c. In the Bode plots

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(Fig. 6b and 6d), a linear region with a slope close to -1 in the frequency range from 1000 Hz to 0.1 Hz can be observed in the impedance moduli, and the corresponding

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phase angles evolve between 60° and 80°, which means the formed passive film is mainly capacitive. The literature proposes different models of equivalent circuits to analyze the EIS

data on passive films formed on stainless steels. After numerous trials, the equivalent circuit presented in Fig. 6e, which considered the bilayer structure: a porous outer layer and a relative compact inner layer, was chosen to fit the EIS spectra obtained in this experiment. CPE (Q, n) is the constant phase element that describes the frequency 12

dispersion behavior caused by the surface heterogeneity of the electrode. The impedance of the CPE (ZCPE) is defined as [10, 34]

Z CPE 

1 Q(  i) n

(1)

and the capacitance (CCPE) is calculated using the following equations by Hirschorn et al. [35, 36]: (2)

g  1  2.88(1  n)2.375

(3)

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CCPE  gQ( d 0 )1n

where Q and n are the parameters of CPE (n = 1, an ideal capacitance; 0.5 < n < 1, a

nonideal capacitance),  is the angular frequency, i is the imaginary number (i2 = -1),  is the dielectric constant (15.6 for passive film of stainless steels), 0 is the vacuum

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permittivity (8.85 × 10-12 F·m-1) and ρd is the resistivity and assigned a value of 500

Ω·m-1 [35]. In this circuit, Rs is the solution resistance, CPE1 represents the

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electrochemical behavior of the outer layer, R1 is the diffusion resistance for ions across the outer layer, CPE2 corresponds to the electrochemical response of the electrical

transfer resistance.

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double layer at the end of the ion channels in the outer layer, and R2 reflects the charge-

The fitting results of the EIS spectra displayed in Fig. 6 are listed in Table 2. The

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fitted R2 values are approximately one order of magnitude larger than the R1 values, which indicates that the charge transfer is the determinant process of the

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electrochemical system of the passive film. Polarization resistance (Rp), which was always used to estimate the stability of the passive film formed on stainless steels, was

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calculated based on

Rp  R1  R2

(4)

where R1 and R2 are the parameters of the fitting procedure. Fig. 7 displays the obtained Rp values of passive films formed on wrought and SLM 316L SSs at different potentials. For wrought 316L SSs, the Rp value increased with the positive shift of film formation potential in the range from -0.10 – 0 VSCE, and decreased with the potential increasing from 0 VSCE to 0.10 VSCE. This result revealed that the passive film formed on the 13

wrought 316L SS became more stable until the formation potential exceeded 0 VSCE, which is in line with the results of the appearance of current fluctuations when polarized at 0.05 and 0.10 VSCE. Unlike the wrought 316L SS, the Rp value of the SLM sample increased gradually as the film formation potential was shifted positively within the selected range, meaning that the stability of the passive film formed on SLM 316L SS increased at more noble potentials. Additionally, the polarization resistance of wrought 316L SS was larger than the SLM sample at low formation potentials, and the converse situation was observed at high formation potentials.

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3.3.2 Mott-Schottky analysis In general, the passive films of stainless steels exhibited a semiconductive

behavior, which can be characterized using Mott-Schottky. Based on the Mott-Schottky theory, a space charge layer exists in the passive film, and there is a linear relationship

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between the inverse square of the pace charge capacitance (1/C-2) and the applied potential (E) neglecting the influence of the Helmholtz layer [23, 29, 34]:

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1 2 T  ( E  EFB  ) for n-type 2 C  0eN D e

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1 2 T  ( E  EFB  ) for p-type 2 C  0eN A e

(5)

(6)

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where NA and ND are the acceptor and donor densities, respectively, EFB is the flat band potential, e is the electron charge (1.602 × 10-19 C),  is Boltzmann constant (1.38 × 1019

J/K), T is the Kelvin temperature (K), and  and 0 have the same meaning as before.

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Fig. 8a and 8b shows the measured Mott-Schottky curves of the passive films formed on the wrought and SLM 316L SSs in 1 M NaCl solution, pH 3 at different film

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formation potentials. A linear region with a positive slope in the applied potential range of -0.2 ~ 0 VSCE was observed in each Mott-Schottky curve. The positive slope reflected that the passive film had the properties of the n-type semiconductor, namely, the predominant carriers of the passive film were oxygen vacancies and/or cation interstitials [37, 38]. Fig. 8c presents the flat band potentials (EFB) of the passive films formed on the wrought and SLM 316L, which were calculated based on the fitting results of the linear regions included in Fig. 8a and 8b. All the EFB values were nearly 14

identical (approximately -0.32 VSCE), and they were approximately 0.15 VSCE smaller than the OCP (not presented here). Fig. 8d displays the donor densities (ND) of the passive films. The calculated ND values were in the range of 1020–1021, mimicking the results reported in previous literature [39]. For the wrought 316L SS, the ND values first decreased with positive shifts of the formation potential and then increased when the film formation potential surpassed 0 VSCE, while the ND values of the passive films formed on SLM 316L SS gradually decreased with the formation potential becoming more positive within the selected range. In addition, the donor density in the passive

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film formed on wrought 316L SS was smaller compared with the SLM sample at the low film formation potential, and the converse situation was observed at high formation potentials. 3.3.3 XPS analysis

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To study the compositions of passive films, XPS measurements were performed

on the wrought and SLM 316L SSs after polarization at -0.1 and 0.1 VSCE for 1 h in 1

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M NaCl solution, pH 3. The XPS results revealed that all the formed passive films had a similar composition of chemical elements, which mainly included Fe, Cr and O. After

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removing the baseline according to Shirley, the XPS spectra of Fe 2p3/2, Cr2p3/2 and O1s were deconvoluted to determine the oxide states of these elements. As shown in

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Fig. 9a, four peaks were contained in the high solution spectra of Fe 2p3/2, i.e., the metallic state Fe0 (707.1 eV), FeO (708.4 eV), Fe2O3 (710.9 eV) and FeOOH (711.8 eV). It is noteworthy that Fe3O4 was difficult to distinguish from FeO and Fe2O3 for the

ur

XPS measurements; thus, Fe3O4 was considered as FeO·Fe2O3, in the compositional percentage calculation [40-42]. Fig. 9b reveals that the element Cr existed in the formed

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passive films in the forms of Cr2O3 (576.1 eV) and Cr(OH)3 (577.4 eV), which are usually considered the main constituents of the inner layer of passive films of stainless steels [34, 43]. Three peaks associated with O2- (530.6 eV), OH- (532.0 eV) and H2O (532.9 eV) constituted the O1s spectrum (Fig. 9c), indicating that bound water was contained in the passive films. The percentages of passive films were obtained based on the fitted peak areas, as shown in Fig. 10. The Feox+hy/Crox+dy ratio clearly revealed that the passive films of 15

wrought 316L SS polarized at -0.1 VSCE was smaller than that polarized at 0.1 VSCE, while for the SLM sample, the passive films formed at -0.1 VSCE had a larger Feox+hy/Crox+hy ratio than the other, potentially because oxide dissolution of Cr is promoted by the occurrence of metastable pitting. In summary, the results of the EIS, Mott-Schottky and XPS measurements suggested that the passive films formed on wrought 316L SSs at low formation potentials were more stable than those formed on the SLM sample, while with increasing film formation potential positivity, the passive film of the wrought sample

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was broken down preceding SLM 316L SS. 3.4 Pit growth of wrought and SLM 316L SSs 3.4.1 Potentiostatic polarization

-p

Fig. 11 illustrates the i-t curves of the wrought and SLM 316L SSs during pit

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growth under potentiostatic polarization. Some differences in the shapes of the i-t curves between the two types of stainless steels could be clearly observed. For the

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wrought 316L SS, the current transients appeared after polarization for a very short time, followed by a sudden increase in the current (Fig. 11a), while for the SLM sample, the current maintained steady in the initial hundreds of seconds and then slowly rose (Fig.

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11b). The differences revealed that the wrought 316L SS was more likely to suffer from metastable pitting corrosion than the SLM specimen. Additionally, the incubation times

ur

of stable pitting corrosion of the two kinds of samples, defined as the polarization time with a current equal to 0.5 mA·cm-2 in this measurement, were markedly different. As

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displayed in Fig. 11a and 11b, the incubation time of wrought 316L SS was 3879 ± 304 s, which was approximately 1550 s longer than the incubation time of the SLM sample (2318 ± 1227 s). This result indicated that metastable pitting corrosion of SLM 316L SSs evolved more easily to stable pitting corrosion than the wrought sample. Fig. 12 shows the morphologies of the two kinds of stainless steels after polarization at 0.25 VSCE in 0.1 M NaCl solution, pH 3. On the wrought 316L SS, hemispherical pits were randomly distributed, and the phenomenon of through-wall 16

penetration could also be observed, as shown in Fig. 12a. Fig. 12b clearly shows that the hemispherical pit always had a smooth interior morphology. The polished pit, representing one of the typical pits observed on stainless steels, reflected that its growth was controlled by the diffusion process during anodic polarization. Fig. 12c reveals that pitting corrosion of SLM 316L SS was initiated at the manufactured defects, among which the gas pores were the most common sites. The preferential dissolution of gas pores is in line with the findings reported by Schindelholz et al., who found using microelectrochemical measurements that the inside wall of the gas pore has a higher

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corrosion tendency than the matrix [17]. According to this rationale, preferential dissolution should also occur on the internal surface of the gas pores with a metal cover,

which was difficult to observe by SEM before the metal cover was completely ruptured,

as presented in Fig. 12d. Additionally, a porous oxide layer clearly covered the bottom

-p

of the gas pore of SLM 316L SS, in contrast to the polished pit of the wrought sample. Generally, the porous oxide layer on stainless steels is caused by electrochemical

re

dissolution, which might occur under conditions in which the passive state is difficult to maintain [44-46]. Therefore, it can be deduced that the pit growth initiated at the gas

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pore of SLM 316L SS occurred mainly in the form of electrochemical dissolution of the passive film under this condition.

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3.4.2 Immersion test

After immersion in 6% FeCl3 solution at 60 ºC for 24 h, the morphologies of the wrought and SLM 316L SSs were obtained. As shown Fig. 13a, hemispherical pits were

ur

randomly distributed on the surface of the wrought sample, and the density of the pits was approximately 1333 cm-2. Additionally, the crystallographic morphologies were

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present at the bottom of the formed pits (Fig. 13b and 13c). As reported by Frankel et al., the crystallographic morphology will be generated during a pit growth process without a salt film for an extended time when immersed at OCP or polarized at a relatively low potential [47]; otherwise, polished pits will form due to salt-filmed dissolution at a relatively high potential [48-50]. Accordingly, it can be seen that pit growth in the immersion test was controlled by active state dissolution of the metal, which is a microstructurally sensitive process [50, 51]. Fig. 13d displays the distribution 17

of the pits formed on the SLM sample, and the pit density was approximately 355 cm2

, which is a quarter of that of the wrought 316L SS. From Fig. 13e – 13i, we can see

that the following two types of pits appeared on the surface of SLM 316L SS after the immersion test: one showed that the pit had no circular striations, which was named Type-I in this paper (Fig. 13e and 13g); the other showed that the pit was characterized by circular striations, which was named Type-II (Fig. 13h and 13i). Type-I pits with the crystalline internal morphology demonstrated that pit growth of SLM 316L was also controlled by active state dissolution as observed for the wrought sample (Fig. 13e and

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13f). Additionally, the appearance of the porous cover confirmed that development under a cover was contained in the pit growth process of the SLM sample like the

extensively reported pit growth framework of stainless steels (Fig. 13g) [50, 52, 53]. The internal wall of Type-II was covered by a porous oxide layer, suggesting that the

-p

pit growth process was controlled by electrochemical dissolution of the passive film

(Fig. 13h and 13i). Generally, the porous oxide layer on stainless steels was caused by

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electrochemical dissolution of the passive film, which usually occurs under conditions in which the passive state is difficult to maintain [44-46].

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To obtain the detailed geometrical features of the pits of the two kinds of stainless steels, CLSM observation was performed. Fig. 14 displays the 3D images and

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corresponding depth profiles of some representative pits generated for the wrought and SLM 316L SSs after immersion for 24 h. The pit shown in Fig. 14a is the largest one appearing on the wrought sample, with a diameter (D) and depth (d) of 213.3 m and

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195.4 m, respectively. The values of D and d of the largest pit found on the SLM 316L SS are larger than those of the wrought sample, reaching up to 422.7 m and 569.3 m,

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respectively (Fig. 14b). This result demonstrated that the SLM 316L SS suffered from more severe pitting corrosion than the wrought one. As can be seen in Fig. 14c and 14d, both the diameter and depth of the Type-I pit were much larger than the manufactured gas pores, while the size of the Type-II pit was nearly equal to the original gas pore. Thus, it could be speculated that the Type-II pit was related to metastable pitting, and the appearance of the Type-I pit resulted from stable pitting. Based on the surface observations, the pitting process of SLM 316L SS could be indicate that pitting 18

corrosion initiated at the gas pore, followed by metastable growth in the form of electrochemical dissolution of the passive film, followed by active dissolution in some metastable pits indicating the transition to stable growth when a certain condition was attained. The sizes of the pits obtained with CLSM were statistically analyzed to provide a visual comparison of the pitting degree between the two type of stainless steels. Fig. 15a shows the coordinate graphs of the randomly selected 100 pits of wrought 316L SS, where most of the points are located around a straight line with a slope of 0.39 in the

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range of (D: 0 – 60 m, d: 0 – 60 m). Fig. 15b shows the frequency distributions of the diameter and depth of the pit, which could be fitted with a logarithmic normal function:

(7)

-p

 1  1  exp   2 (ln x   ) 2  , x  0  f ( x,  ,  )   2 x  2  0, x  0 

re

where the function f(x, , ) is the frequency, x represents the pit diameter (D) and depth (d) here, and the parameters of  and  are the mean value and variance, respectively.

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The calculated mean values of D and d are 10.8 m and 6.4 m, respectively. In addition, the ratio of D/2d is always used to characterize the shape of the pits: the pit with D/2d = 1 is the normative hemispheric, D/2d < 1 is related to a deep pit, and D/2d > 1 reflects

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a shallow pit [54, 55]. As shown in Fig. 15c, the frequency of the D/2d value was also in line with the logarithmic normal distribution, and the fitted mean value was 1.21,

ur

which revealed that the pits of wrought 316L had a shallow hemispherical shape. For SLM 316L SS, only five Type-I pits were found on the entire surface of the tested SLM

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sample. The five Type-I pits are plotted in Fig. 15d, and the distributions of their diameters and depths are presented in Fig. 15e; their mean diameter and depth were 166.1 m and 155.9 m, respectively. Similarly, the values of D/2d for the five pits were also calculated. As shown in Fig. 15f, their mean D/2d value was 0.85, which is smaller than the wrought sample and suggests a deep hemispherical shape. Statistical analysis of Type-II pits (manufactured gas pores) of SLM 316L SS was also conducted for comparison. The diameters, depths and D/2d values of the Type-II pits clearly were 19

more discretely distributed than the pits presented on the wrought 316L SS, as shown in Fig. 15g – 15i. The mean D/2d values of the Type-II pits were 1.08, indicating that this type of pit had a nearly hemispherical shape. From the statistical results and surface observations, we could conclude that a smaller number of pits (including metastable and stable pits) was generated on SLM 316L SS, but the sizes of stable pits (Type-I) was larger, compared to the wrought sample. Thus, the wrought 316L SS had a larger frequency of pitting nucleation and metastable pitting, while the transition from metastability to stability pits was more likely to occur on the SLM sample, and the

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growth rate of the formed stable pit was larger. A supplemented immersion test using an SLM specimen with some marked gas

pores was performed to determine the geometric factors influencing the transformation

from metastable pitting to stable pitting. In the supplemented measurement, the

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immersion time was extended to 48 h to generate more Type-I pits on the immersed

sample. Fig. 16a and 16a’ show the five marked gas pores in the SLM 316L SS, among

re

which the ones marked with “1” and “2” are the pores without the metal cover, and the other three are covered. From Fig. 16b and 16b’, it is apparent that the Type-I pits

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formed at the sties marked with “3”, “4” and “5”, and no significant changes occurred at the other two gas pores during the immersion test for 48 h. This result revealed that

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the metastable pitting initiating at covered gas pores was easier to transfer to stable pitting than that induced by uncovered gas pores.

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4. Discussion

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4.1 Pitting mechanism induced by gas pores In general, the pitting corrosion of stainless steels comprises three distinct stages,

namely, nucleation, metastable growth and stable growth, where the nucleation and transition from metastable to stable pit growth are the two critical events that affect the pitting process [52, 56]. The nucleation of pitting corrosion related to the breakdown of the passive film always occurs at the defects (such as inclusions, boundaries, 20

dislocations and so on) where the worse passive film is formed [40]. The transition from metastability to stability pit requires that the environment inside the pit can maintain the continuous active dissolution state at the pit surface [57]. The experimental results in this manuscript confirm that pitting corrosion of SLM 316L SS initiates at gas pores, and the metastable pits at pores with metal covers are easier to transfer into stable pits than open pores (Fig. 13 and 16). Fig. 17 displays the schematic diagram of the pitting mechanism induced by the manufactured gas pores of SLM 316L SS. In this framework, the gas pore is considered

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a spherical cap, and r, R and h are the cross-sectional radium, spherical radium and depth of the spherical cap, respectively, as shown in Fig. 17a. When h < R, the spherical cap represents an open gas pore (Fig. 2a and 2b); otherwise, it denotes a gas pore with

a metal cover (Fig. 2c and 2d). The relationship of the three geometrical parameters of

-p

a spherical cap is defined as

re

r 2  2Rh  h2

(8).

1 V   Rh2   h3 3 S  2 Rh

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and its volume (V), surface area (S) can be calculated as

(9) (10).

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The aggressive pit environment required by pitting corrosion mainly constitutes H+ and Cl-, where H+ is derived from hydrolysis of metal cations and Cl- migrates into

ur

the pit to balance the charge and maintain electroneutrality [47]. Hence, the accumulation of metal cations can be regarded as the cause of the harsh local conditions,

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and the concentration of cations at the pit surface (Csurf) are used to assess the harshness of the electrolytes inside the pits. For SLM 316L SS, the accumulation of cations inside the gas pore is related to the electrochemical reactions at the pit surface and the transportation of metal cations out from the pit interior into the bulk solution (Fig. 17b). In the initial stage, the gas pore environment is not aggressive, and the passive film can be formed on gas pore surface. It is well known that passive films of stainless steels are in dynamic equilibrium between film growth and dissolution, among which film 21

dissolution is the electrochemical reaction accompanied by the generation of metal cations [40]. Therefore, the electrochemical dissolution of passive film should be the generation process of cations in the gas pore in this stage. According to Moffat et al. [58], the film dissolution rate (id) is dependent on the potential drop (φf/s) as follows:

id  id0 exp(

 mF f / s RT

)

(11)

where α represents the polarizability of the film/solution interface (α = 0.7 [59, 60]), and id0 and m are the exchange current density and the number of electrons.

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Simultaneously, the diffusion of metal cations driven by the concentration difference occurs at the rate (idiff) defined by [52, 57]:

idiff 

mFDCsurf deff

(12)

-p

where F is the Faraday constant, D is the effective diffusivity of metal cations, and deff

is the effective diffusion length that is dependent on geometrical parameters (R, h and

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r). Compared to the matrix, gas pores have a larger film dissolution rate and a smaller diffusion rate, which is beneficial for the accumulation of metal cations. Thus, the

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beginning of accumulation of metal cations/electrochemical dissolution of passive film in the gas pore can be regarded as the nucleation of pitting corrosion.

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The metastable pitting stage occurs following nucleation (Fig. 17c). In this stage, the pH value of the gas pore environment drops with the accumulation of metal cations, which results in a decreasing passive film thickness [23, 31]. The thinning of the passive

ur

film could influence both the film dissolution rate and cation diffusion rate. For the given applied potential (Eapp), if the change in carrier density of the passive film is

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ignored, the field intensity (E0) should be a constant, and the film dissolution rate can be written as id  id0 exp(

 n f / s RT

) exp(

 n f / s RT

)  id0 exp(

 n f / s RT

) exp(

 nE0Csurf RTK d

)

(13)

where Kd is the change in thickness of the passive film per the cation concentration. The film dissolution rate will clearly be enhanced by the accumulation of metal cations. In addition, the thickness of the passive film is always at the nanolevel that is 3 – 4 22

orders of magnitudes smaller than the size of the gas pores, so the variation in the geometrical parameters of the spherical cap can be neglected in this stage. Additionally, the diffusion rate of the metal cation can still be described using Eq. 12, which also increases with cation accumulation like the film dissolution. In this stage, for id > idiff, the accumulation rate of metal cations in the gas pore is d Csurf / dt  id  idiff  id0 exp(

 n f / s RT

) exp(

 nE0Csurf RTK d

)

nFDCsurf deff

(14)

 n f / s  nE0  ln c1  id0 exp( ) RT RTK d  Csurf (t )   nE0 / RTK d

 t   c exp( nFDt ) 2 deff

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and the relationship between the cation concentration and reaction time can be obtained:

(15)

where c1 and c2 are the constant of integration. When id = idiff, the steady state is obtained,

-p

and the cation concentration reaches a maximum value (Csurf, max) that is independent of the time.

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With further accumulation of metal cations and decreases in pH values in the gas pore, the cation concentration is likely to reach the critical value (Ccrit), where the

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thickness of the passive film would decrease to 0, and active state dissolution of metal could occur instead of electrochemical dissolution of the passive film at the gas pore surface. Therefore, the accumulation of metal cations in this stage is dependent on the

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active state dissolution of the matrix and transport of metal cations. The active dissolution rate (iad) can be written as (neglecting the total ohmic potential drop of bulk

ur

and pit solutions) [47]

Eapp  Ecorr

a

(16)

)

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iad  icorr exp(

where icorr and Ecorr are the corrosion current and potential of the bare metal, and a is the anodic Tafel slope. Therefore, the active dissolution rate could be considered as a constant for the given applied potential (Eapp) and temperature (T), as reported by Frankel and Burstein [47, 52]. Additionally, it is noteworthy that the influence of the active dissolution on the geometrical parameters of the gas pore is sufficiently large to be considered. As shown in Fig. 13 – 15, the active state dissolution promotes the 23

transfer of the gas pore into the hemispheric pit with a metal cover (for h > R) or without a metal cover (for h  R), as shown in Fig. 17d and 17e. The diffusion rate in the hemispheric pit during the stage can be written as follows [47, 57, 61]:

idiff 

3nFDCsurf 3nFDCsurf  for open pit: 2 R ' 2 ( R  KV iad t )

(17)

idiff 

crnFDCsurf crnFDCsurf  for covered pit: 2 R'  ( R  KV iad t )2

(18)

where R’ is the pit radium that is related to the time, Kv is the volume of metal dissolved

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per unit anodic charge, c is a dimensionless constant with a value of 2 – 3, and the mouth radium is replaced using the cross-sectional radium (r) here for simplicity. According to the framework of pitting corrosion proposed by Frankel et al., when the

cation concentration in the pit can maintain iad > idiff, stable pit growth would take place.

-p

Hence, stable pitting corrosion will occur at the open gas pore (h  R) when Csurf > Ccrit. For the covered gas pore (h > R), active dissolution could occur at least until the cover

re

ruptures as the critical metal concentration (Ccrit) is achieved. After the cover is completely ruptured, if the metal concentration can still support the active dissolution,

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the metastable pit will transfer to stable growth, and otherwise it will be repassivated. The relationship between the cation concentration and reaction time can be obtained

na

when Csurf > Ccrit:

Csurf (t )  Ccrit  iad (t  tcrit )  [ R  KV iad (t  tcrit )]3nFD /2 KV iad for open pit

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Csurf (t )  Ccrit  iad (t  tcrit )  exp{

crnFD } for covered pit  KV iad [ R  KV iad (t  tcrit )]

(19) (20)

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where tcrit is the time for the critical cation concentration. As the accumulation process continues, the metal cations in the formed pit are likely to be saturated (Csat), and the salt film will be deposited on the pit surface. Therefore, the active state dissolution of the metal is controlled by the diffusion process in the solution saturated with metal cations, which is distinguished from the active dissolution controlled by the charge transfer process when Csat > Csurf > Ccrit.

24

4.2 Influence of passive film properties on pitting nucleation From the pitting mechanism introduced above, pitting nucleation of SLM 316L SS corresponds to the electrochemical dissolution of the passive film formed on the gas pore surface, which is dependent on the carrier density and/or the space charge layer of the passive film. Fig. 18a presents a schematic diagram of the space charge layer formed on the passive film/electrolyte interface according to the Mott-Schottky measurement and the energy band theory of the semiconductor: The flat band potential is the electrode potential where the band is flat, namely, the

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i.

Femi level of the semiconductor is equal to that of the redox couple in the

electrolyte [62, 63]. Therefore, the flat band potential measured by Mott-Schottky should correspond to the equilibrium potential of the redox couples in the solution;

In this work, the measured flat band potential was close to the corrosion potential

-p

ii.

and more negative than passive region, as displayed in Figs. 3 and 9, meaning that

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the Femi level of the redox couple was more positive than that of passive film, and the electrons always tended to be transferred from a high Femi level to a low Femi

lP

level based on the above contents;

iii. The Femi level of the n-type semiconductor was close to its value band, with a greater number of carrier density resulting in closer positioning.

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In the schematic diagram, the accumulation layer forms and the energy band bends downward, accompanied by electrons migrating from the electrolyte to the passive film,

ur

and ΦSC represents the electric voltage of the space charge layer. By definition, we could conclude that the smaller donor density would result in a larger Femi level

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difference between the passive film and the redox couples in solution. This phenomenon will drive the transfer of more electrons from the solution to the passive film, thus resulting in a larger ΦSC. If the Helmholtz layer is ignored, ΦSC is the potential drop at the film/solution interface (φf/s), as displayed in Fig. 18b. Accordingly, as for the passive film with a larger donor density, the electrochemical dissolution of the passive film in the gas pore (Eq. 11) is more rapid. The enrichment of dislocations at the surface of the gas pore results in passive film formation here with a larger donor 25

density than the matrix [64, 65]. Therefore, the generation rate of metal cations in the gas pore is easier than the matrix. 4.3 Influence of geometrical parameters on pit growth Based on the pitting mechanism induced by the gas pores of SLM 316L, the pitting corrosion was influenced by the geometric parameters of the gas pore only through changing the cation diffusion rate, because both the electrochemical dissolution rate of the passive film and active dissolution rate of the metal were independent of the gas

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pore shape. In Eq. 12, the effective diffusion length (deff) is associated with the shape

of the gas pore. For simplicity, the effective diffusion length is roughly considered the ratio of the volume to cross-section area of the spherical cap:

V h h3    r 2 2 6r 2

(21)

-p

deff 

6r 2 nFDCsurf 3r 2 h  h3

(22)

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idiff 

re

Combining Eq. 12 and Eq. 21

Obviously, with a larger r and smaller h, the deff value decreases and idiff increases. To further demonstrate the influence of the geometric parameters, the variation of

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electrochemical dissolution (id or iad), cation diffusion rate (idiff), cation concentration (Csurf) in the gas pore (pit) and radium (R or R’) of the gas pore (pit) were systematically

ur

analyzed. According to Eq. 13 – 22, it could be speculated that the following three cases might occur during the pitting corrosion initiated at gas pores with different geometric

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parameters.

For Case-I, the gas pore with a smaller deff value and the increase rate of idiff with

time were relatively large, and the steady state (id = idiff) would form at point “F” during electrochemical dissolution of the passive film (AB), as shown in Fig. 19a. At this point, the cation concentration reached a maximum value (Cmax) that was smaller than the critical concentration (Ccrit) for the active state dissolution of metal (Fig. 19b). Because no active dissolution occurred in this case, the size of the gas pore showed no significant 26

change (Fig. 19c). In the experimental results of this paper, the pitting corrosion displayed in Fig. 13h and 13i could be categorized as this case, where the pitting process initiated at open gas pores was halted at the metastable growth stage, and the bottoms of the gas pores were covered with a porous oxide layer caused by electrochemical dissolution of the passive film. As for Case-II, with an increasing deff value, the idiff increase rate dropped, which would result in variation curves of id (AB) and idiff (OE) that do not intersect before the critical value (point B) is attained, as shown in Fig. 19d. After point B, active

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dissolution of the metal occurred at a constant rate (iad), suggesting pit growth transitions from metastability to stability. With the accumulation of metal cations in the

formed pit, the diffusion rate continued to increase until the steady state, where idiff was equal to iad (BF). Then, the cation concentration continued to increase with the

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accumulation process in the steady state (FC), as shown in Fig. 19e. At point C, the solution in the pit was saturated by metal cations. Subsequently, a salt film was

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deposited at the pit surface, and the controlled process of pit growth was transformed from charge transfer (BC) to diffusion (CD). Fig. 19f plots the variation in pit size in

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this case, where the pit grows at a constant rate in the stage (BC) controlled by charge transfer, and the growth rate decreases with time in the stage (CD) under diffusion

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control.

Concerning Case-III, with a continually increasing deff value, the solution in the gas pore tended to be saturated (point C), as presented in Fig. 19g. In this case, the

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variation tendencies of the cation concentration and pit size were similar to those observed in Case-II, as shown in Fig. 19h and 19i. When comparing Fig. 19d and 19g,

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it could be clearly determined that the differences between the generation rate (id/iad) and diffusion rate (idiff) in stages AB and BC of Case-III were larger than those in the same stages of Case-II. This phenomenon resulted in a shorter time requirement for the critical metal concentration (tcrit) and saturation (tsat) for Case-III compared with CaseII, namely, the pitting corrosion induced by the gas pores in Case-III was easier to transform from metastable growth to stability. Based on the results of the immersion test, pitting, which grows under change transfer control as displayed in Fig. 13e, should 27

stop at stage BC of Case-II/Case-III, and the pits shown in Fig. 13g possessed features of stage CD of Case-II/Case-II, where the pit growth was controlled by the diffusion process. In a summary, the gas pores with a smaller r and larger h exhibited a larger diffusion resistance and were more likely to induce stable pitting corrosion. This finding is in line with the experimental results shown in Fig. 16, where stable pitting corrosion occurred at covered rather than open gas pores.

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4.4 Influence of the harshness of the bulk solution on pitting corrosion Based on the pitting mechanism introduced in Section 4.1, the pitting processes

induced by the gas pores of SLM 316L SS were determined by the accumulation of metal cations in the gas pores, i.e., the electrochemical dissolution of the passive film

-p

(id) and diffusion of cations (idiff). Therefore, it is necessary to discuss the influence of

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pH values and Cl- content in the bulk solution on the pitting resistance of SLM 316L SS from two aspects.

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Electrochemical dissolution of passive film (id)—As shown in Eq. 13, the pH values of the gas pore environment could influence the electrochemical dissolution rate of passive film by decreasing its thickness and increasing the potential drop at the film/solution

na

interface. With the decreasing pH value of the bulk solution, the original thickness of the passive film was reduced. Thus, the initial dissolution rate of the passive film

ur

increased, namely, the variation curves of id were shifted to left. Enhancing the Clcontent in bulk solution could increase the donor density of the formed passive film [37,

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66]. As discussed in Section 4.2, a larger donor density in the passive film results in a larger potential drop at the film/solution interface. Therefore, increasing the Cl- content could also cause the variation curves of id to shift to the left. As shown in Fig. 20a, after the left shift of id curves, they were more likely to reach the critical concentration (Ccrit) for active state dissolution, and the required time (tcrit) became shorter. Diffusion of metal cations (idiff)—Apart from the dissolution of passive film, both the pH value and Cl- content of the bulk solution could also influence the transport process 28

of metal cations. As we introduced before, the aggressiveness of the gas pore environment was mainly composed of H+ and Cl- and could be assessed based on the cation concentration (Csurf). Likewise, the harshness of the bulk solution was estimated using a hypothetical cation concentration (Cbulk). Therefore, considering the original pH value and Cl- content, the diffusion rate of cations in the gas pore (Eq. 14) could be rewritten as

idiff 

nFDCsurf nFDCbulk  deff deff

(23)

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where the aggressive ions diffuse from the gas pore to the bulk solution only when Cbulk < Csurf. This phenomenon demonstrates that the diffusion curve related to Eq. 23 is shifted to the right with respect to the curve reflected by Eq. 12. From Fig. 20b, it is apparent that the right shift of the diffusion curve would make the critical value for

-p

active dissolution of metal easier to obtain. Summarily, increasing the aggressiveness

of the bulk solution (with a smaller pH value and larger Cl- content) could promote the

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pitting corrosion induced by the gas pore of SLM 316L SS through the left shift of the

(Fig. 3 and Fig. 4).

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dissolution curve of the passive film and the right shift of the diffusion of metal cations

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4.5 Pitting mechanism of wrought 316L SS

Generally, the pitting corrosion of wrought 316L SS was initiated at MnS

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inclusions. The pitting process at MnS inclusions could be concluded from previous reports [67-70]. First, MnS inclusions were dissolved in chloride-containing solutions,

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which caused the formation of thiosulfate ions and weak acidification, according to Eq. 24 [69, 70]:

2MnS + H2O  2Mn 2  S2O32  6H  8e

(24).

In the next step, elemental sulfur was deposited on and around the inclusions though the disproportionation reaction shown in Eq. 25 and the reduction reaction in Eq. 26 [68]:

S2O32  2H  S  SO2  H2O

(25) 29

S2O32  6H  4e  2S  3H2O

(26)

The coexistence of element sulfur and chloride ions led to dissolution of the steel matrix, followed by the formation of trenches at the inclusion/steel matrix interfaces [71]. Then, the hydrolysis reaction of metal cations released from the steel matrix dissolution were accompanied by a decreasing pH value and concentration of chloride ions in the local environment, which would result in the transition from the passive to the active state in the bottom trenches. Once the chloride concentrations and pH values inside the trenches exceed the critical values for autocatalytic growth, stable pitting corrosion occurs [71].

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The size and geometric shape of the MnS inclusions played important roles in their

capacity to form stable pitting corrosion. As reported by Krawiec et al., larger MnS inclusions start to dissolve before small ones [72], and the dissolution of small

inclusions is difficult to provide sufficient aggressive ions for pit self-propagation [73].

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Ruoru et al. found that only the MnS inclusions with sizes larger than 1 μm could induce

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the pitting corrosion of 304 stainless steel in chloride-containing solutions [74]. The geometric shape of the MnS inclusions could influence the diffusion process of

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aggressive ions inside the formed trenches. In general, it was easier to initiate pitting corrosion of stainless steel for the deep MnS inclusions than shallow ones because the shallow inclusions exhibited high diffusion rates and thus were effective for the

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dispersal of aggressive ions in the formed trenches [72, 75]. In our previous work, we confirmed that pitting corrosion of stainless steels was

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more likely to occur at MnS inclusions than oxide inclusions [76]. In this work, the spherical inclusions of SLM 316L SS had sizes in the range from 50 – 200 nm,

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indicating that the inclusions in SLM 316L SS had little susceptibility to pitting corrosion. Therefore, the pitting corrosion initiated at the inclusions was only considered in the wrought 316L SS. 4.6 Different pitting behavior of wrought and SLM 316L SSs The results of the immersion test confirmed that the frequency of metastable pitting of SLM 316L SS was lower, and the transition from metastability to stability 30

pits was easier, compared to the wrought sample. The difference in pitting corrosion between the two types of stainless steels could be attributed to their different pitting mechanisms. From the perspective of pitting nucleation, the pitting corrosion of the wrought 316L always initiated at MnS inclusions, the density and chemical activity of which were larger than the gas pores for which pitting corrosion nucleated the SLM sample. Metastable pit growth following the preferential dissolution of MnS inclusions occurred in the mode of active state dissolution of the metal [53], the current density of which was much larger than that of electrochemical dissolution of the passive film (Fig.

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19a). Therefore, the frequency and current density of the metastable pitting of the wrought 316L SS was much larger than the SLM specimen. Concerning the transition

from metastability to stability pits, the pits formed on the wrought 316L SS had nearly

hemispheric shapes with average sizes that were smaller than those of the gas pores

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(Fig.13b and 15). According to Eq. 12, it can be clearly discerned that the cation diffusion rate was larger in the pits formed on wrought 316L SS than the manufactured

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gas pores of SLM 316L SS. Hence, the requirement for stable pit growth (iad = idiff [47]) for SLM 316L SS could be reached more easily than the wrought 316L SS, namely, the

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metastable pit growth of SLM 316L SS was more likely to transform to stable growth. Additionally, the potentiodynamic polarization curve results revealed that SLM 316L

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SS exhibited superior pitting corrosion resistance in solution with slight aggressiveness, and less pitting corrosion resistance in media with strong aggressiveness, compared to the wrought 316L SS. This phenomenon could be due to the different control processes

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of the pitting corrosion of the wrought and SLM 316L SSs in solution with different aggressiveness. When the two types of stainless steels were exposed to an extremely

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aggressive condition, their pitting corrosion was controlled by the transition from metastability to stability pits. Due to the easier transition, SLM 316L SS was more sensitive to pitting corrosion than the wrought sample in aggressive solution (1 M NaCl, pH 1 and 3 M NaCl, pH 3). With a decreasing aggressiveness of the conditions, the pitting corrosion of SLM 316L SS would be transferred to the process controlled by pitting nucleation related to passive film breakdown preceding the wrought one, because of the more stable passive film of SLM 316L (Fig. 7 and 8). Therefore, it could 31

be deduced that SLM 316L SS possessed superior pitting corrosion resistance to the wrought one in less aggressive solutions (0.01 M – 1 M NaCl, pH 3; 1 M NaCl, pH 3 – 11). With further decreases in aggressiveness, the pitting corrosion of the two types of stainless steels were under control of the nucleation process, and both exhibited good pitting resistance under conditions with much less aggressiveness (1 M NaCl, pH 13).

5. Conclusion The pitting corrosion of the wrought and SLM 316L SSs was comparatively

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investigated in solutions with different aggressiveness in the current paper. The main conclusions are as follows.

The potentiodynamic polarization curves revealed that the SLM 316L SS exhibited

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higher sensitivity to pitting corrosion in extremely aggressive solutions (1 M NaCl, pH

1 and 3 M NaCl, pH 3), while under other conditions, the wrought sample was more

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vulnerable to pitting.

According to the results of the EIS and Mott-Schottky measurements, the passive

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films formed on wrought 316L SSs at low formation potentials were more stable than the SLM sample, but with more positive film formation potential, the passive film of the wrought sample was broken down prior to SLM 316L SS.

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The immersion test showed that the wrought 316L SS had a higher frequency of metastable pits; however, the metastable pits generated on SLM 316L SS transitioned

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more easily to stable growth.

Based on the in-situ experiment, it was confirmed that the metastable pits initiated

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at covered gas pores of SLM 316L SS could transition to stable growing pits more easily than those initiated at open gas pores.

Competing Interest Statement

All the Authors (Zhiwei Duan, Cheng Man, Chaofang Dong, Zhongyu Cui, Decheng Kong, Li Wang, Xin Wang) declare no Competing Financial or Non-Financial Interests. 32

Author statement

We deeply appreciate your consideration of our manuscript. No conflict of interest exits in the submission of this manuscript, and manuscript is approved by all authors for publication. I would like to declare on behalf of my co-authors that the work described was original research that has not been published previously. Author statement: Zhiwei Duan: Conceptualization, Methodology, Data Curation.

Chaofang Dong: Writing - Review & Editing. Zhongyu Cui: Data curation, Writing - Review & Editing.

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Decheng Kong: Software, Investigation.

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Cheng Man: Validation, Supervision, Writing - Original Draft.

Li Wang: Visualization, Data Curation.

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Xin Wang: Project administration.

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Acknowledgements

This work is supported by the National Key Research and Development Program of China (No. 2017YFB0702300), National Natural Science Foundation of China (No.

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51901216, 51871028), the Fundamental Research Funds for the Central Universities (No. FRF-TP-17-002B), and China Postdoctoral Science Foundation (No.

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2019M652471).

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Figure list

Fig. 1. The microstructures of wrought and SLM 316L SSs: (a) and (b) IFP of wrought

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and SLM 316L SS, where the gray lines and dark lines are used to identify the low angle grain boundary and high angle grain boundary, respectively; (c) and (d) SEM

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images of SLM 316L SS; (e) and (f) TEM images of SLM 316L SS and EDS results of the marked inclusion; (g) and (h) SEM images of wrought 316L SS and EDS results of

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the marked inclusion.

Fig. 2. The gas pores with different morphologies in SLM 316L SS: (a) and (b) the hemispherical gas-pores without covered metal; (c) and (d) gas-pores covered by

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remnant metal.

Fig. 3. The cyclic potentiodynamic polarization curves of wrought and SLM 316L SSs

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in different solutions: (a) and (b) wrought and (b) 316L SSs in solutions containing different NaCl contents (pH = 3); (c) and (d) wrought and SLM 316L SS in 1M NaCl solutions at different pH values. Fig. 4. The pitting behavior of the wrought and SLM 316L SSs: (a) and (b) potentiodynamic polarization curves in 1M NaCl at pH = 3; (c) pitting potentials in solutions with different NaCl contents at pH = 3; (d) pitting potentials in 1 M NaCl solutions with different pH values. 41

Fig. 5. i-t curves of wrought and SLM 316L SSs during the passive film formation: (a) and (c) the curves in linear and double logarithmic coordinates for SLM 316L SS; (b) and (d) the curves in linear and double logarithmic coordinates for wrought 316L SS. Fig. 6. The EIS results of the wrought and SLM 316L SSs: (a) and (b) Nyquist and Bode of the SLM 316L SS; (c) and (d) Nyquist and Bode of the wrought 316L SS; (e) the equivalent circuit used in this work. Fig. 7. The polarization resistance of the passive film formed on the wrought and SLM 316L SSs polarized at different anodic potentials in 1 M NaCl solution at pH = 3.

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Fig. 8. The Mott-Schottky of the wrought and SLM 316L SSs: (a) wrought 316L SS, (b) SLM 316L SS, (c) the EFB and OCP of the two samples, and (d) the fitted ND of the two samples.

Fig. 9. The detailed XPS spectra of Fe 2p3/2 (a), Cr 2p3/2 (b) and O 1s (c) of the passive

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films formed on wrought and SLM 316L SSs through potentiostatic polarization at ±0.10 VSCE in the solution with 1 M NaCl, pH = 3.

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Fig. 10. The constitute percentages and Feox+hy/Crox+hy ratios of the passive films formed on wrought and SLM 316L SSs through potentiostatic polarization at ±0.10 VSCE in the

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solution with 1 M NaCl, pH = 3.

Fig. 11. i-t curves of the wrought and SLM 316L SSs during the pit growth under

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potentiostatic polarization in the solution with 1 M NaCl, pH = 3: (a) wrought 316L SS; (b) SLM 316L SS.

Fig. 12. The morphologies of wrought and SLM 316L SSs after potentiostatic

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polarization at 0.25 VSCE in 0.1 M NaCl solution with pH = 3: (a) and (b) wrought 316L SS; (c) and (d) SLM 316L SS.

Jo

Fig. 13. The morphologies of the wrought and SLM 316L SSs after immersed in 6% FeCl3 solution at 60 ºC for 24 h: (a) – (c) wrought 316L SS; (d) – (i) SLM 316L SS. Fig. 14. The typical pitting 3D morphologies and depth of the wrought and SLM 316L SSs after immersed in 6% FeCl3 solution at 60 ºC for 24 h: (a) wrought 316L SS; (a) – (c) SLM 316L SS.

42

Fig. 15. The statistical analysis of the pits generated on the wrought and SLM 316L SSs: (a) – (c) wrought 316L SS, (d) – (f) Type-I pits of SLM 316L SS, and (g) – (i) Type-II pits of SLM 316L SS. Fig. 16. The 2D and 3D morphologies of the SLM 316L SS before and after supplemented immersion test: (a) and (a’) before immersion test, (b) and (b’) after immersion test. Fig. 17. The schematic diagram of the pitting mechanism induced by the manufactured gas-pores of SLM 316L SS.

ro of

Fig. 18. (a) The schematic diagram of the space charge layer formed at the passive film/electrolyte interface, (b) the schematic diagram of potential distribution.

Fig. 19. The variation curves of the idiff, id, Csurf and R in the gas-pore with different

geometrical parameters: (a), (d) and (g) for idiff and id; (b), (e) and (h) for Csurf; (c), (f)

-p

and (i) for R.

Jo

ur

na

lP

the right shift of the idiff curve.

re

Fig. 20. The variation of the id and idiff curves: (a) the left shift of the id curve, and (b)

43

ro of -p re lP na ur Jo Fig. 1. The microstructures of wrought and SLM 316L SSs: (a) and (b) IFP of wrought and SLM 316L SS, where the gray lines and dark lines are used to identify the low angle grain boundary and high angle grain boundary, respectively; (c) and (d) SEM images of SLM 316L SS; (e) and (f) TEM images of SLM 316L SS and EDS results of 44

the marked inclusion; (g) and (h) SEM images of wrought 316L SS and EDS results of

Jo

ur

na

lP

re

-p

ro of

the marked inclusion.

45

ro of -p re

Fig. 2. The gas pores with different morphologies in SLM 316L SS: (a) and (b) the

lP

hemispherical gas-pores without covered metal; (c) and (d) gas-pores covered by

Jo

ur

na

remnant metal.

46

ro of -p

re

Fig. 3. The cyclic potentiodynamic polarization curves of wrought and SLM 316L SSs in different solutions: (a) and (b) wrought and (b) 316L SSs in solutions containing

lP

different NaCl contents (pH = 3); (c) and (d) wrought and SLM 316L SS in 1M NaCl

Jo

ur

na

solutions at different pH values.

47

ro of -p re

Fig. 4. The pitting behavior of the wrought and SLM 316L SSs: (a) and (b)

lP

potentiodynamic polarization curves in 1M NaCl at pH = 3; (c) pitting potentials in solutions with different NaCl contents at pH = 3; (d) pitting potentials in 1 M NaCl

Jo

ur

na

solutions with different pH values.

48

ro of -p

re

Fig. 5. i-t curves of wrought and SLM 316L SSs during the passive film formation: (a) and (c) the curves in linear and double logarithmic coordinates for SLM 316L SS; (b)

Jo

ur

na

lP

and (d) the curves in linear and double logarithmic coordinates for wrought 316L SS.

49

ro of -p re lP na

ur

Fig. 6. The EIS results of the wrought and SLM 316L SSs: (a) and (b) Nyquist and Bode of the SLM 316L SS; (c) and (d) Nyquist and Bode of the wrought 316L SS; (e) the

Jo

equivalent circuit used in this work.

50

ro of

Fig. 7. The polarization resistance of the passive film formed on the wrought and SLM

Jo

ur

na

lP

re

-p

316L SSs polarized at different anodic potentials in 1 M NaCl solution at pH = 3.

51

ro of -p

re

Fig. 8. The Mott-Schottky of the wrought and SLM 316L SSs: (a) wrought 316L SS, (b) SLM 316L SS, (c) the EFB and OCP of the two samples, and (d) the fitted ND of the

Jo

ur

na

lP

two samples.

52

Fig. 9. The detailed XPS spectra of Fe 2p3/2 (a), Cr 2p3/2 (b) and O 1s (c) of the passive

Jo

ur

na

lP

re

-p

±0.10 VSCE in the solution with 1 M NaCl, pH = 3.

ro of

films formed on wrought and SLM 316L SSs through potentiostatic polarization at

53

ro of

-p

Fig. 10. The constitute percentages and Feox+hy/Crox+hy ratios of the passive films formed

on wrought and SLM 316L SSs through potentiostatic polarization at ±0.10 VSCE in the

Jo

ur

na

lP

re

solution with 1 M NaCl, pH = 3.

54

Fig. 11. i-t curves of the wrought and SLM 316L SSs during the pit growth under

ro of

potentiostatic polarization in the solution with 1 M NaCl, pH = 3: (a) wrought 316L SS;

Jo

ur

na

lP

re

-p

(b) SLM 316L SS.

55

ro of -p re

Fig. 12. The morphologies of wrought and SLM 316L SSs after potentiostatic

lP

polarization at 0.25 VSCE in 0.1 M NaCl solution with pH = 3: (a) and (b) wrought 316L

Jo

ur

na

SS; (c) and (d) SLM 316L SS.

56

ro of -p

re

Fig. 13. The morphologies of the wrought and SLM 316L SSs after immersed in 6%

Jo

ur

na

lP

FeCl3 solution at 60 ºC for 24 h: (a) – (c) wrought 316L SS; (d) – (i) SLM 316L SS.

57

ro of

Fig. 14. The typical pitting 3D morphologies and depth of the wrought and SLM 316L

-p

SSs after immersed in 6% FeCl3 solution at 60 ºC for 24 h: (a) wrought 316L SS; (a) –

Jo

ur

na

lP

re

(c) SLM 316L SS.

58

ro of -p

Fig. 15. The statistical analysis of the pits generated on the wrought and SLM 316L SSs:

re

(a) – (c) wrought 316L SS, (d) – (f) Type-I pits of SLM 316L SS, and (g) – (i) Type-II

Jo

ur

na

lP

pits of SLM 316L SS.

59

ro of

Fig. 16. The 2D and 3D morphologies of the SLM 316L SS before and after supplemented immersion test: (a) and (a’) before immersion test, (b) and (b’) after

Jo

ur

na

lP

re

-p

immersion test.

60

ro of

Fig. 17. The schematic diagram of the pitting mechanism induced by the manufactured

Jo

ur

na

lP

re

-p

gas-pores of SLM 316L SS.

61

ro of

Fig. 18. (a) The schematic diagram of the space charge layer formed at the passive

Jo

ur

na

lP

re

-p

film/electrolyte interface, (b) the schematic diagram of potential distribution.

62

ro of -p re lP

Fig. 19. The variation curves of the idiff, id, Csurf and R in the gas-pore with different

Jo

ur

and (i) for R.

na

geometrical parameters: (a), (d) and (g) for idiff and id; (b), (e) and (h) for Csurf; (c), (f)

63

ro of

Fig. 20. The variation of the id and idiff curves: (a) the left shift of the id curve, and (b) the right shift of the idiff curve.

Table 1 Typical fitted electrochemical parameters for EIS of wrought and SLM 316L

Q1 Rs

(10

potential

Ω·cm-

Ω-

(VSCE)

2

n1

1

·cm-

316L SS

Jo

0

-0.05

2

)

n2

1

·cm-

Ω·cm2

)

2

·sn)

lP

(10

5

4.40

0.91

5.31

10.01

3.39

0.84

7.75

±

±

±

±

± 4.35

±

±

±

0.42

0.13

0.06

1.42

0.53

0.12

1.46

na

SLM

Ω-

R2

5.32

5.26

3.19

0.98

6.69

22.94

3.70

0.87

9.37

±

±

±

±

± 2.64

±

±

±

0.33

0.68

0.01

0.87

0.63

0.09

1.63

4.97

2.04

0.91

5.44

5.56

0.81

15.41

±

±

±

±

±

±

±

0.51

0.75

0.05

1.21

1.32

0.11

2.10

5.22

2.28

0.92

5.14

2.87

0.86

11.43

±

±

±

±

±

±

±

0.43

0.31

0.03

0.31

0.32

0.06

3.01

ur

0.05

Ω·cm)

·sn)

(10-5

(μFcm-

2

2

0.10

(10

C1

4

re

Applied

Q2

R1

-5

-p

SS in 1 M NaCl, pH = 3 at different apllied potentials

4.67 ± 1.33

6.14 ± 0.42

64

C2 (μFcm2

)

2.51 ± 0.98

Rp (105 Ω·cm2

)

8.28 ± 2.88

4.43 ± 1.21

10.01 ± 2.50

2.56 ± 0.83

16.03 ± 3.31

2.93 ± 0.81

11.92 ± 3.32

Wrought 316L SS

0

-0.05

3.64

0.84

±

±

±

±

±

±

0.35

0.90

0.02

1.33

0.93

0.08

4.96

5.44

0.96

4.05

28.17

1.78

0.83

19.15

±

±

±

±

± 4.91

±

±

±

0.52

1.01

0.02

0.69

0.31

0.12

3.12

5.05

4.90

0.97

3.34

29.90

2.05

0.86

12.23

±

±

±

±

± 8.21

±

±

±

0.61

0.52

0.06

0.97

0.57

0.06

1.36

4.66

2.12

0.99

2.87

17.98

2.25

0.87

8.59

±

±

±

±

± 7.42

±

0.23

0.62

0.01

1.56

5.24

3.67

0.95

2.27

16.13

6.93

±

±

±

±

± 6.32

±

0.49

0.91

0.04

0.79

5.25

3.40

0.92

4.63

±

±

±

0.58

0.89

0.02

6.87 ± 2.34

0.86

9.16 ±

±

1.42

0.62

Jo

ur

na

lP

-0.10

3.21

65

5.07± 0.95

2.69 ± 0.63

5.39 ± 2.28

1.13 ± 0.61

19.19 ± 3.81

2.09 ± 0.93

12.57 ± 2.33

2.70 ±

8.86

ro of

0.05

0.91

±

±

0.04

1.85

0.86

6.52

±

±

-p

0.10

3.00

0.96

0.03

0.95

2.08

0.86

3.21

±

±

±

0.39

0.05

0.67

re

-0.10

5.20

0.74

±

3.41

7.08 ± 2.31

6.75 ±

1.74 2.12 ± 0.95

3.68 ± 1.39