Improvement of corrosion resistance of SS316L manufactured by selective laser melting through subcritical annealing

Journal Pre-proof Improvement of corrosion resistance of SS316L manufactured by selective laser melting through subcritical annealing Chengshuang Zhou, Shiyin Hu, Qiaoying Shi, Huimin Tao, Yangyang Song, Jinyang Zheng, Peng Xu, Lin Zhang

PII:

S0010-938X(19)31273-9

DOI:

https://doi.org/10.1016/j.corsci.2019.108353

Reference:

CS 108353

To appear in:

Corrosion Science

Received Date:

20 June 2019

Revised Date:

14 October 2019

Accepted Date:

17 November 2019

Please cite this article as: Zhou C, Hu S, Shi Q, Tao H, Song Y, Zheng J, Xu P, Zhang L, Improvement of corrosion resistance of SS316L manufactured by selective laser melting through subcritical annealing, Corrosion Science (2019), doi: https://doi.org/10.1016/j.corsci.2019.108353

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Improvement of corrosion resistance of SS316L manufactured by selective laser melting through subcritical annealing

Chengshuang Zhoua, Shiyin Hua, Qiaoying Shia, Huimin Taoa, Yangyang Songa, Jinyang Zhengb,*, Peng Xuc, Lin Zhanga,*

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a College of Materials Science and Engineering, Zhejiang University of Technology, Hangzhou 310014, China b institute of Process Equipment, Zhejiang University, Hangzhou 310027, China c Hefei General Machinery Research Institute, Hefei 230031

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Graphical abstract

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Highlights:

 Subcritical heat-treatment greatly improves pitting corrosion resistance.

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 Pitting initiation of stainless steel by SLM is primarily at melt pool boundary.

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 Corrosion improvement is related to microstructure evolution.  950℃ heat-treatment can dissolve the melt pool boundary.  The corrosion of SLM 316L mainly depends on interface and dislocation.

ABSTRACT The effect of heat treatment on pitting-corrosion resistance of 316L stainless steel manufactured via selective laser melting (SLM) was investigated. The result shows that both pitting-corrosion resistance and mechanical properties under Subcritical-temperature heat treatment (950 °C) are better than 1100 °C and as-received. For as-received samples, the most prone to pitting corrosion is melt-pool boundary (MPB). Subcritical-temperature heat treatment not only eliminates MPBs and high-density dislocation but also retains advantages of inclusion nanocrystallisation,

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cellular substructure. And recrystallisation occurs during 1100 °C heat treatment significantly

reduces low-angle grain-boundary density and increases segregation and re-precipitation, which weakens corrosion resistance and mechanical properties.

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Keywords: A. stainless steel; B. polarisation; B. polarization; B. SEM; B. TEM; C.

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pitting corrosion

1. Introduction

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Selective laser melting (SLM) is an additive-manufacturing technology that melts powdered raw material using a high-energy laser beam before its solidification. It has been frequently adopted.

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The SLM additive-manufacturing method offers the advantages of material savings, higher

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efficiency, capability for manufacturing of complicated components, and suitability for small batch production compared with the traditional wrought iron method [1-3]. In the past decade, additive manufacturing tremendously grew. Numerous studies have shown that the application of SLM not only produces complex alloy parts that cannot be manufactured by conventional methods but also exhibits superior mechanical properties compared with other casting methods [4-6].

The 316L austenitic stainless steel (SS) is widely used in aerospace, automotive, marine, medical equipment, and other fields owing to its low price, excellent mechanical properties, and corrosion resistance [7-9]. The 316L SS made by SLM (SLM SS316L) are currently produced in large quantities. Many previous studies demonstrated that SLM SS316L shows higher yield strength and ductility than conventional wrought SS. Although the SLM manufacturing has made significant progress in the mechanical properties of alloys [10, 11], it is still in its infancy stage in terms of the

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corrosion properties and corrosion mechanisms for alloys, and further research is needed [12].

Sun et al. [13] demonstrated that the general corrosion behaviour of SLM SS316L is similar to that of wrought SS316L under a 0.9 wt.% NaCl solution. However, the SLM SS is more susceptible

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to pitting-corrosion attacks because of its lower breakdown potential, which is attributed to its

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porosity results. In contrast with the result by Sun et al., Sander et al. [14] found that the pitting potential of SLM SS316L samples is higher than that of the wrought SS316L samples. Their theory

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stated that inclusions were refined or annihilated during the SLM manufacturing process. Thus, the

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number of pit nucleation sites in the SLM SS316L samples was reduced. They also ruled out the effect of porosity on the pitting corrosion because the porosity was between 0.01% and 0.4% in

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their work. Chao et al. [15] revealed that the SLM manufacturing process avoided MnS inclusion formation and Cr-depletion zone. Instead, it formed nanoscale oxide inclusions, which improved the

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pitting-corrosion resistance. The studies of Sander [14] and Chao [15] only considered the influence of inclusions on the pitting corrosion. The influence of interfacial factors [especially the unique melt-pool boundary (MPB) and grain boundaries (GBs)] as important defects in crystals was not considered. Trelewicz et al. [16] showed that the formation of solute segregation around the cellular structures was detrimental to the pitting-corrosion resistance of SLM SS316L. However, Gorsse et

al. [17] confirmed that no significant Cr and Mo contents were depleted at the cellular-structure boundaries. We need to note that no experimental data existed to prove where the pitting-corrosion initiation preferentially occurs, and no post observation of the microstructure at pit initiation sites has been made. Suryawanshi et al. [18] found that the corrosion resistance of the SLM SS316L samples after 500 °C heat treatment was higher than that of wrought SS316L because of the refinement of the structure. However, the reason why the pitting-corrosion resistance of the

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high-temperature heat treatment was different from that of the as-received SLM samples was not explained. Kong et al. [19] found that the anti-corrosion property and durability could be improved using the fraction of the oxides and thickening of the passive film after heat treatment. However, the

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microstructure evolution after the heat treatment and its effect on corrosion were not explained in

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their article.

In many actual production processes, immediately eliminating the residual stress using post

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heat treatment following the additive-manufacturing component production is routine [20, 21].

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From the previous work done by the aforementioned researchers, we can observe that heat treatment influences the corrosion resistance. Although many corrosion studies have been reported,

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a number of related issues remain unaddressed, e.g. pitting-initiation location, correlated heat treatment, microstructure, and corrosion property.

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In the present article, these issues are discussed in detail. In this work, potentiodynamic and

potentiostatic polarisations were carried out to evaluate the corrosion resistance of SLM SS316L. The surface morphology and microstructure were characterised using an optical microscope (OM), scanning electron microscopy (SEM), electron backscattered diffraction (EBSD), and transmission electron microscopy (TEM).

2. Experimental 2.1.

Sample and heat-treatment preparation

SLM SS316L samples, which were manufactured using an EOS 250 machine (Germany), were used. The chemical compositions of both traditional wrought SS316L and SLM SS316L are listed in Table 1, which indicates that the Cr and Mo content in SLM SS316L was less than that of the

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wrought SS, and the other elements were basically the same. Two heat-treatment conditions were applied to the SLM SS316L samples: (1) HT950—the temperature increased to 950 °C from room temperature and was kept at 950 °C for 4 h in vacuum and subsequently furnace-cooled. (2)

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HT1100—the temperature increased to 1100 °C from room temperature and kept at 1100 °C for 1 h

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in vacuum followed by furnace-cooling. By using a differential scanning calorimetry test, we found

complete recrystallisation.

Electrochemical measurements

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2.2.

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that the recrystallisation temperature of SLM SS316L was 1035.2 °C. HT2 was selected to generate

The samples used in the electrochemical measurements were embedded in a plastic tubing with

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epoxy resin; this was accomplished after the rear of the samples were welded with a wire. The

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working area was a 10-mm-diameter circle exposed in a solution. All samples were sequentially ground using from 400 to 2000 grit sandpaper, polished using 0.5- and 0.1-μm diamond polishing agent, rinsed with ethanol, cleaned using deionised water, and finally dried in air. Electrochemical measurements, such as potentiodynamic and potentiostatic polarisations, were performed using an Ivium SXRe electrochemical workstation. We adopted the use of a conventional three-electrode cell in which the sample acted as the working electrode, platinum was the counter

electrode, and saturated calomel electrode was the reference electrode. Electrochemical tests were performed in a 3.5 wt.% NaCl solution. All electrochemical measurements were carried out in a thermostatic water bath at 25 °C. Five replicate tests were performed under each condition. An intergranular corrosion test was performed in 10% oxalic acid solution. The samples acted as the anode, and the SS sheet acted as the cathode. The current density was 1 A/cm2, and the samples were etched for 90 s at 25 °C. All samples were washed with running water, dried in air,

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and finally observed using a microscope. The solution that corrodes the precipitate phase is 7.5g FeCl3 + 25mL HCl + 50mL H2O solution. After 60s of corrosion, rinse with water and dry in air.

Microstructure characterisation

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2.3.

OM was used to observe the metallographic phase and pit-initiation sites. SEM was used to

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observe the microstructure morphology and pit-initiation sites. The grain size, GBs, and

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crystal-orientation distribution were investigated by EBSD experiments. In the EBSD experiments, the acceleration voltage, collection speed, and scanning step were 20 kV, 637.76 Hz, and 3 μm,

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respectively. TEM was used to observe the distribution of the dislocations and sub-GBs. The mercury intrusion porosimetry (MIP) and Brunauer,Emmett,Teller’s test (BET) method were

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3. Results

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adopted to measure the porosity.

3.1.

Effect of heat treatment on microstructure evolution

The metallographic phase of the as-received, HT950, and HT1100 samples is shown in Fig. 1. Fig. 1(a) shows the fan-shaped melt pools as-received samples with an average size of approximately 100 μm and with irregular fine grains. These fan-shaped melt pools are typical of an

SLM sample microstructure and are product of incremental layer-by-layer construction process in the SLM process [22]. After the 950 °C heat treatment, the MPB was dissolved, and the grain morphology did not significantly change [Fig. 1(b)]. After the 1100 °C heat treatment, the grain morphology significantly changed from irregular to regular grains which was similar to the shape of the traditional wrought SS316L. This result proved that recrystallisation occurred at 1100 °C, and the grain shape was transformed [Fig. 1(c)].

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The grain morphologies and grain sizes of the as-received, HT950, and HT1100 samples were compared using EBSD, as shown in Fig. 2. We can see that the grain shape of the as-received samples from the inverse pole figure (IPF) [Fig. 2(a)] was irregular and lathy, whereas no

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significant difference in the grain shape was observed in the HT950 samples compared with the

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as-received samples. The grain shape of the HT950 samples was also curved and lathy, as shown in Fig. 2(b). However, the grains of the HT1100 samples obviously became more regular polygonal

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grains, as shown in Fig. 2(c), which was attributed to the recrystallisation.

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Fig. 3 shows the quantitative description of the grain shape and grain size. The average grain diameters of the as-received, HT950, and HT1100 samples were 28.11, 28.74, and 36.02 μm,

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respectively, which indicated that a minor change in the average grain diameter of the HT950 samples appeared compared with the as-received samples. However, the average diameter directly

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increased after the HT1100 treatment. The grain-aspect ratios exhibited a downward trend of 3.32, 3.25, and 2.22, respectively, indicating that the HT1100 grains tended to be equiaxed, which was attributed to the recrystallisation at 1100 °C. The results show that the distributed grain shape and grain size were more equiaxed and homogeneous after the HT1100 treatment, which may be beneficial to the grain rotation and migration [20]. For the relationship between the grain size and

corrosion resistance, previous research has shown that the grain size exerted no significant effect on the pitting corrosion of austenitic SS [23]. Fig. 4 shows the GB distribution of the as-received, HT950 and HT1100 samples. The black line represents the high-angle GBs (HAGBs; misorientation of more than 15°), and the green line represents the low-angle GBs (LAGBs; misorientation of less than 15°). The red column represents the distribution of the proportion of the misorientation angle. Based on the misorientation-angle

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distribution bar graph, we can observe that the HAGB distribution was uniform but it also had a relatively low fraction of the as-received samples. Correspondingly, the LAGB fraction was

relatively high. The misorientation-angle distribution of the HT950 samples was similar to that of

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the as-received samples, but the LAGB fraction was slightly reduced compared with that of the

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as-received samples. Compared with the as-received and HT950 samples, the HAGB fraction in the HT1100 samples significantly increased. Meanwhile, the calculated average misorientation angles

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for the as-received sample, HT950, and HT1100 were 13.83°, 15.02°, and 23.72°, respectively [24].

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The increase in the average misorientation angle indicated that the LAGBs were dissolved and migrated to the HAGBs after the HT1100 treatment [25]. We could also observe in the

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grain-boundary distribution images that the LAGB fractions of the HT950 samples barely changed when compared with those of the as-received samples; however, the fractions of the HT1100

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samples were significantly reduced. A mechanical test of the three SLM SS316L samples was performed, and the result is listed in

Table 2, which indicates that the yield stress was reduced with the increase in the heat-treatment temperature. The elongation values of the HT950 and HT1100 samples were similar and were both higher than that of the as-received samples. HT950 demonstrated the highest tensile strength than

the other samples. In summary, HT950 exhibited excellent mechanical property.

3.2.

Effect of heat treatment on corrosion properties

The potentiodynamic polarisation curves of the as-received, HT950, HT1100 samples as well as that of traditional wrought SS316L in 3.5 wt.% NaCl at 25 °C are shown in Fig. 5(a), which shows that as-received SLM SS316L had a lower pitting potential (Epit) than traditional wrought

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SS316L. This result may be attributed to the low Cr and Mo content in SLM SS316L [26]. The pitting potential of HT1100 was similar to that of wrought SS316L, and Epit of the HT950 sample was significantly higher than those in the other conditions. According to the potentiodynamic

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polarisation curves, a range fluctuation occurred in the transpassivation zone in the as-received

curve. However, the curves did not show obvious fluctuation in the transpassivation zone after heat

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treatments of 950 and 1100°C. The cathodic branches in the potentiodynamic polarisation curves

process of the SLM samples [27].

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almost overlapped; this indicated that the microstructure had almost no effect on the cathodic

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The electrochemical parameters obtained from the potentiodynamic polarisation curves are listed in Table 3. The low corrosion current density (icorr) exhibited a slow corrosion rate, and the

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large difference between the pitting potential (Epit) and corrosion potential (Ecorr), denoted as ΔE,

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indicated a high corrosion resistance [28]. Their icorr values were almost the same, but ΔE of HT950 was much larger than those in the other conditions, as listed in Table 3, which confirmed that HT950 has excellent pitting-corrosion resistance compared with the others. To further evaluate the corrosion resistance under the environment of 3.5 wt.% NaCl solution, the potentiostatic polarisation (0.05VSCE) was measured. The potentiostatic polarisation curves are shown in Fig. 5(b). The current density of the samples rapidly dropped due to the rapid passivation

of the surface at the start. The current densities of the HT950 and HT1100 samples were basically stable, whereas that of the as-received samples greatly fluctuated, which may be related to the more surface defects of the as-received samples. All curves were stable after 15000 s, and the current-density sequence was as follows: as-received > HT1100 > HT950, indicating that the corrosion-resistance order was HT950 > HT1100 > as-received, which was consistent with the passive current density (ipass) from the potentiodynamic polarisation curves listed in Table 3.

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To evaluate the position where pitting preferentially occurred, we observed that it stopped at the point where the pitting potential had just passed though the transpassivation zone when the

potentiodynamic polarisation test was performed. The location of the trigger pits is shown in Fig. 6,

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which shows that the pitting-initiation sites of the as-received samples were at the MPBs [Fig. 6(a)].

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Pitting was triggered at the GBs because of MPBs were dissolved in the HT950 samples [Fig. 6(b)]. The pitting sites also mainly sprouted at the GBs in the HT1100 samples [Fig. 6(c)]. To more

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accurately understand the location of the pitting-initiation sites, multiple samples were counted. We

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used a 1000× magnification OM to observe the pits, which was sufficient to distinguish between the pitting pits and pores and made the statistical work easier because of the higher contrast and wide

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field of view of the OM. A typical feature of the pits was regular circular pits with smooth inner walls and a size range from 8 to 12 µm. A total of 343 pits were counted, including 157 pits in the

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as-received samples, 81 pits in the HT950 samples, and 105 pits in the HT1100 samples. The statistical results are shown in Fig. 6(d). We found that the number of pits that appeared in the MPBs was the largest in the as-received samples, followed by that in the GBs. The smallest was in the grains. For the HT950 samples, pitting initiation mainly occurred at the GBs, and the proportion of pitting inside the grains was very low. In the HT1100 samples, pitting initiation also mainly

occurred in the GBs, but the pitting proportion inside the grains was significantly higher than that in the as-received and HT950 samples. According to the statistical results, pitting initiation most likely occurred at the MPB then at the GB. Pitting initiation first occurred at the GB when the MPBs were dissolved. After the heat treatment of 1100 ℃, in addition to the GBs where pitting initiation mainly occurred, the pitting initiation inside the grains also increased. To further confirm the position where pitting preferentially occurred, intergranular corrosion

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experiments were performed. The surface morphologies of the as-received, HT950, and HT1100 samples after intergranular corrosion are shown in Fig. 7. Fig. 7(a) shows that the MPBs were the main attack position for corrosion in the as-received samples. The corrosion grooves at the MPBs in

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the as-received samples were quite clear and coarse, but those at the GBs were fairly fuzzy. After

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the high-temperature heat treatments of 950 and 1100 ℃, the MPBs were dissolved, and the GBs became the primary corrosion position. The distribution of the corrosion sites in the HT1100

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samples was more disorderly than that of the as-received and HT950 samples. Not only was serious

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corrosion present at the GBs but also a considerable number of corrosion pits were distributed in the grains [Fig. 7(c)]. This result was consistent with the results shown in Fig. 6(d). Compared with the

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HT1100 samples, the corrosion positions of the HT950 samples almost occurred only at the GBs [Fig. 7(b)]. The corrosion grooves were blurry and fine, demonstrating that the degree of

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intergranular corrosion in the HT950 samples was weaker than that of the others. The results of the intergranular corrosion again indicated that the corrosion preferentially attacked the MPBs when MPBs existed, and the corrosion would mainly attack the GBs when the MPBs were dissolved. Many corrosion sites appeared in the grains except for the corrosion sites at the GBs with the HT1100 heat treatment. This result indicated that recrystallisation not only caused the change in the

GB to make the intergranular corrosion serious but also may have caused the inclusion in the grains to re-precipitate, leading to many corrosion sites in the grains.

4. Discussion The previous reports [15, 19, 29] indicated that some inclusions appeared in SS316L, manufactured via SLM, including macro-inclusions and nano-inclusions. Although pitting etch pits

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were observed using SEM, we found that the pits were smooth and flat inside without other

corrosion products or inclusions in the as-received [Fig. 8(a)], HT950 [Fig. 8(b)], or HT1100 [Fig. 8(c)] samples. We can conclude that no macro-inclusion appeared in the pitting initiation zone. The

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rapid heating and cooling rates during the SLM process affected the particle-size distribution and

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concentration of the inclusions; this caused the inclusions to become very small or annihilated to contribute in triggering the pits [14], which led to no obvious segregation around the GBs. This

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result indicated that the primary cause of the pitting initiation was not the micro-inclusions.

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However, we also did not exclude the possibility that nano-inclusions may have caused the pitting corrosion. Nano-inclusions still existed in SLM SS316L [30]. Previous reports [15, 27, 30]

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indicated that circular nano-inclusions with a size of 20–200 nm existed and were mainly composed of Si, Cr, and O and may also contain elements such as Mn and Al. Saeidi et al. [30] found that

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depletion of Si and Cr occurred around the circular nano-inclusions, which confirmed that the formation of nano-inclusions absorbed two elements from the neighbouring area; this caused the pitting corrosion to be more easily initiated at the nano-inclusions. However, the nano-inclusions may in fact have been taken off after the pitting corrosion. It was difficult to observe whether the pitting corrosion occurred at the position of the inclusions through post observation [31]. We could

only reasonably believe that the pitting corrosion was more probable to occur at the inclusion sites. Porosity is also one of the important factors that affect the corrosion resistance of SLM SS316L. Although pores are difficult to avoid, it can be controlled to extremely low levels by optimising the production process [27]. Sun et al. [32] suggested that porosity would have a preferential pit initiation if it was more than 1%, whereas the porosity of the SLM SS316L samples used in the present experiment was approximately 0.9%, as measured by the BET and MIP method,

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which was below 1%. Zhang et al. [33] also confirmed that porosity was not necessarily the reason for the reduction in the pitting potential; thus, it had almost no effect on the pitting potential in the present work.

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The highest pitting potential was obtained when the grain orientation was in the most

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closed-packed [101] direction compared with the others [34]. However, we could see from the IPF images that no obvious difference existed in the grain orientation. In addition, the ratio of the [101]

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direction in the as-received sample was substantially consistent with those of the HT950 and

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HT1100 samples, as observed from the calculations. The result showed that no direct relationship existed between the grain orientation and corrosion resistance of the different heat-treated SLM

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SS316L samples. Based on the aforementioned discussions, we believe that the primary factors that led to the reduction in the pitting resistance and pit initiation of SLM SS316L were the interface,

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MPBs, and GBs. Therefore, the microstructure evolution was related to the corrosion resistance of SLM SS316L.

The morphologies of the cellular sub-structures are shown in Fig. 9. Numerous sub-grains existed inside the grains of the SLM SS316L samples, which was consistent with previous research results [35, 36]. Figs. 9(a) and (b) show that a significant change in the microstructure can be found

between the as-received and HT950 samples. The as-received samples demonstrated a cellular dislocation wall structure with an average size of approximately 1 μm [37]. After a 950 °C heat treatment, the cellular dislocation structure was dissolved, and a spherical cellular trace microstructure appeared at the position of the original cellular structure; this indicated that the dislocation density decreased [38]. The size of the cellular structure also changed, and the range was below 0.5 μm. A small proportion of the spherical structure remained from the HT1100 heat

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treatment in which the average size was similar to that of the HT950. However, majority of the

sub-grains disappeared into grains without a sub-grain structure, such as that shown in the lower left corner in Fig. 9(c).

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To compare the microstructures, bright-field TEM images are shown in Fig. 10. A sub-grain

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was clearly observed in the as-received grains, as shown in Fig. 10(a). The dislocations in the cellular structure were dense and chaotic. The cellular-structure boundaries were entangled by a

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mass of dislocations, which formed a significant dislocation-rich region at the boundary. We

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believed that the high concentration of dislocations at the cellular sub-GBs was due to the rapid solidification of the austenite grains during the SLM process [39]. In the SLM-production process,

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the dislocation density of the as-received samples was much higher than that of the traditional wrought samples because of the large amount of residual stress generated during the rapid cooling

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process [40]. Although high concentration of dislocations could lead to higher tensile strength and hardness [41], the dislocation-rich regions were also the areas of severe corrosion [42, 43]. The TEM images of the samples after the 950 ℃ heat treatment are shown in Figs. 10(b) and (d). Fig. 10(b) shows that the high temperature caused the chaotic dislocations of the original cellular sub-grains to disappear. On the other hand, a small number of dislocation lines appeared to be

staggered, and the dislocation density was greatly reduced. This result was consistent with the changes in the cellular sub-grains shown in Fig. 9. Some sub-GBs are clearly shown in Fig. 10(d). We can clearly see that the sub-GBs were entangled by the ordered dislocations shown in Fig. 10(c), indicating that some of the dislocations at the sub-GBs migrated and disappeared under the high-temperature condition of the HT950 heat treatment. Meanwhile, the dislocation lines in some sub-GBs almost dissolved and completely formed random GBs, which may have been due to the

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recrystallisation of the partial grains, as shown in the black-arrow position in Fig. 10(d). The

dislocation density in the grains was further reduced, and a large number of dislocations were

dissolved in the sub-grain boundaries after a heat treatment of 1100 ℃. Meanwhile, the GBs became

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regular and straight because of the recrystallisation. However, entangled dislocation lines remained

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in some sub-grain boundaries [19, 44, 45].

The MPBs of the SLM SS316L samples included the heat-affected zone (HAZ) and remelted

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zone. The cellular microstructure in the HAZ was replaced by a granular microstructure, and the

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cellular microstructure in the remelted zone was characterised by coarse grains. The change indicated that the microstructure in the MPBs was inhomogeneous [46, 47]. The local temperature

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gradient caused by the rapid solidification difference at the solid–liquid interface could generate a surface-tension gradient in the MPBs during the SLM manufacturing process [48] and resulted in

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local heat accumulation at the MPBs, which caused the MPBs to have higher interfacial free-energy than that of the matrix and GB. The stress distribution in the vicinity of the MPBs was uneven [46]. Previous studies have found that no obvious element-segregation attributed to the extremely high cooling rate existed after the SLM (typically ~107 K/s) [49] or only a slight element segregation, which included C, O, and Si, occurred at the MPBs [50], which formed a non-metallic unstable state

[46]. However, the slight element segregation may not have been the main reason for the pitting to primarily occur in the MPBs. Meanwhile, the pores were always located along the MPBs even if the porosity was very low, and the stress distribution in the vicinity of the MPBs was not uniform. In summary, the defects, which included the inhomogeneous microstructure, pores, residual stress, high energy, HAZ, and unstable non-metallic elements located in the MPBs, made the MPBs the weakest part in the SLM SS316L samples. Hence, pitting primarily occurred at the MPBs in the

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as-received samples. The pitting potential values of the HT950 and HT1100 samples were higher than that of the as-received samples, which was mainly attributed to the dissolution of the MPBs. The MPBs, which were most vulnerable to pitting attack, were dissolved. When this happened,

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the GBs became the primary location and could develop into a local disturbance. The atomic

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arrangement was disordered at the GB, which was an important crystal defect. The free interfacial energy was higher than that of the substrate, which was beneficial to ion chemisorption to accelerate

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the SLM corrosion. This is also the reason why pitting initiated at the GBs without the MPBs. The

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stainless steel manufactured by the SLM process exhibited the characteristics of rapid cooling. Thus, it offered the advantage of finer inclusion [19]. However, the GB was reformed, and the SLM

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SS316L grains were similar to the wrought SS grains after recrystallisation at 1100 °C. Chromium depletion, which was caused due to chromium carbide, was re-precipitated in the random GBs; this

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led the intergranular corrosion to preferentially attack the random GBs [51]. The quantitative analysis of the precipitated phase of the SLM sample is shown in Fig. 11. The precipitates and inclusions were refined and eliminated because of the rapid cooling of the SLM process, and the element distribution of the SLM samples was more uniform after the 950 °C heat treatment. Thus, no obvious precipitated-phase distribution in the grains was observed [Figs. 11(a) and (b)], which led to

the high corrosion resistance [44]. After the 1100 °C heat treatment, we could clearly observe that many circular precipitates were distributed in the grains with a diameter of approximately 500 nm, as shown in the red-circle positions in Figs. 11(c) and (d); these were attributed to the re-precipitation of the intragranular precipitated phase due to recrystallisation after the 1100 °C heat treatment. Based on the EDS linear-scanning analysis of the intragranular precipitated phase, we found that the precipitated phase was primarily composed of Si, O, and Cr, as shown in Fig. 12. We confirmed that

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compared with the subcritical temperature heat treatment, the chemical segregation of the SLM samples demonstrated more grains after recrystallisation. The experimental results were consistent with the intergranular corrosion results and pitting-corrosion site statistics. Hence, the decrease in the

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pitting resistance after the 1100 °C heat treatment compared with the subcritical temperature heat

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treatment was attributed to the significant increase in the chemical segregation within the grains after recrystallisation.

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In addition to the interface and GB segregation, the cellular substructure in the microstructure

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of the SLM samples was also an important factor that affected the corrosion resistance and mechanical properties. The cellular sub-grain structure was a unique structure in the additive

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manufacturing process. It had a high dislocation density in the cellular sub-grains owing to its process characteristics and rapid cooling [37, 36]. We found that the dislocation density of the

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cellular sub-GBs and interior was significantly reduced, as shown in Fig. 10, and the cellular dislocation wall became a low-energy GB. The increase in the low-energy GB improved the corrosion resistance. Therefore, the corrosion resistance of the samples after a high-temperature heat treatment increased compared with that of the as-received samples. Recrystallisation, GB reconstruction, and LAGB dissolution occurred after the 1100 °C heat treatment. Meanwhile, we

can clearly see from the red histogram shown in Fig. 4 that the high-density LAGBs of the HT950 samples demonstrated no significant change compared with that of the as-received samples. However, the HAGB proportion of the HT1100 samples increased compared with that of the as-received and HT950 samples. This result also proved the GB reconstruction. Recrystallisation would result in inclusion precipitation and elemental segregation, which were similar to the solidification treatment of the wrought SS316L, leading to a decrease in the corrosion resistance

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[52], which was one of the decisive reasons for the lower corrosion resistance of HT1100 than

HT950. The decrease in the dislocation density explained the reduction in the compressive strength of the HT950 and HT1100 samples. These mechanical results were consistent with the

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microstructure-evolution process. The dissolution of the dislocations is also proven in Fig. 10. The

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decrease in the dislocation density also indicated the reduction in the intragranular defects. Further, Figs. 9 and 10 show that the migration and rearrangement of the dislocation made the sub-grain

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structure more stable and more evenly distributed. Thus, the passivation zone was longer, and the

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transpassivation zone of the HT950 and HT1100 samples did not distinctly fluctuate compared with that of the as-received samples.

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As a summary, the MPBs, chemical segregation, proportion of LAGBs, change in the sub-grain structure, and dislocation density were all factors that affected the corrosion resistance of

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SLM SS316L after a high-temperature heat treatment. Moreover, the dissolution of the MPBs (which were the primary attack location of pitting), reduction in the dislocation density in the sub-grains, homogenisation of the non-equilibrium phase, and chemical composition (which retained the advantages of weak GB segregation and inclusion nanocrystallisation during the SLM process) caused the corrosion resistance of the HT950 samples to be significantly higher than that

of the as-received samples. After the high-temperature heat treatment reached 1100 °C, recrystallisation occurred, which also offered advantages to the MPB dissolution; the dislocation density was also reduced. However, the re-precipitation of the precipitated phase within the grains after recrystallisation and the significant increase in the proportion of high-energy GBs offset some of the corrosion resistance; this resulted in worse corrosion resistance than the subcritical temperature samples. However, the corrosion resistance of HT1100 was still higher than that of the

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as-received samples; this indicated that the most critical causes of the corrosion resistance were the dissolution of the MPBs and reduction in the dislocation density in the sub-grains as compared with the other factors.

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This study provided new ideas for the improvement of corrosion resistance and mechanical

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properties of additive-manufacturing alloys. The effect of subcritical temperature heat treatment on

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the stress corrosion will be studied in the future.

5. Conclusion

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In summary, this work has systematically studied the effects of microstructure evolution after different temperature heat treatments of SLM SS316L on the corrosion resistance and

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pitting-initiation position. The main conclusions are presented as follows.

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(1) Subcritical heat treatment of 950 °C (HT950) induced the transformation of cellular substructures from high-energy high-density dislocation wall to a low energy sub-grain boundary, whereas it did not change the LAGBs. In contrast, a heat treatment of 1100 °C (HT1100) resulted in recrystallisation, which greatly reduced the LAGB density. (2) The order of corrosion resistance was HT950 > HT1100 > as-received sample; herein, the

subcritical temperature heat treatment demonstrated the best corrosion resistance. In the as-received samples, the most critical position of the pitting initiation was at the MPBs, whereas the GBs became the primary pitting-initiation site in the HT950 specimens when the MPBs were dissolved. Subsequently, the pitting-initiation site in the grains increased when heat treatment of 1100 °C was applied due to the re-precipitation. (3) The subcritical temperature heat treatment resulted in the best corrosion resistance and

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mechanical properties. The main reason was that the subcritical temperature heat treatment not only retained the advantages of inclusion nanocrystallisation and weak GB segregation during the SLM process but also eliminated the MPBs and high-density dislocation. The

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increase in the precipitated-phase re-precipitation and high-energy GB proportion during

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Conflicts of interest

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was weaker than that of HT950.

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recrystallisation at 1100 °C were the main reasons why the corrosion resistance of HT1100

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There are no conflicts to declare.

Acknowledgements This research was supported by the National Key Basic Research Program of China (973 Program, Grant No. 2015CB057601), the National Natural Science Foundation of China (51571181) and the Zhejiang Provincial Natural Science Foundation of China (LY19E010006).

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Figure captions

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Fig. 1 Metallographic phase images of the (a) as-received, (b) HT950, and (c) HT1100 samples

Fig. 2 IPF of the SLM SS316L obtained though EBSD: (a) as-received; (b) HT950; and (c) HT1100

Fig. 3 Variations in the average grain size and grain-aspect ratio in the SLM 316L SS with various

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sample types

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Fig. 4 EBSD maps of the GB distribution and misorientation-angle distribution bar graph of the (a)

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as-received, (b) HT950, and (c) HT1100 samples

Fig. 5 (a) Potentiodynamic polarisation curves. (b) Potentiostatic polarisation curves of the as-received, HT950, HT1100, and wrought SS316L in 3.5 wt.% NaCl at 25 °C

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Fig. 6 Metallographic phase and pits of the (a) as-received, (b) HT950, and (c) HT1100 samples. (d)

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Statistical chart of the number of pit sites

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Fig. 7 Surface morphologies after intergranular corrosion: (a) as-received; (b) HT950; (c) HT1100

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Fig. 8 SEM images of the pits: (a) as-received; (b) HT950; (c) HT1100

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Fig. 9 SEM images of the SLM SS316L: (a) as-received; (b) HT950; (c) HT1100

Fig. 10 TEM images of the SLM SS316L: (a) as-received; (b)–(d) HT950

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Fig.11 SEM images of precipitated phase (a), (b) HT950 samples, (c)-(d) HT1100 samples

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Fig.12 SEM-EDS line scanning of the precipitate in HT1100 sample

Tables Table 1 Chemical composition of traditional wrought 316L stainless steel and SLM produced 316L stainless steel Material

Ni

Cr

Mo

Mn

Si

P

S

Fe

wrought 10.64 17.05 2.54 1.21 0.43 0.028 0.031 balance 10.53 16.13 2.52 0.23 0.57 0.014 0.006 balance

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SLM

Table 2

σb/MPa

as-received

591.29

669.30

HT950

452.69

HT1100

370.82

δ/%

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σs /MPa

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The mechanic test date of as-received and HT950 samples

36.98 39.22

580.16

40.02

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700.03

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Table 3

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Electrochemical parameters of traditional wrought and SLM samples (Ecorr: pitting potential, Ecorr: corrosion potential, icorr: corrosion current density, ipass: passive current density) As-received HT950 HT1100 Wrought

Epit(vSCE) 0.122±0.002 0.794±0.010 0.337±0.004 0.425±0.005

Ecorr(VSCE) -0.3133±0.0027 -0.3668±0.0005 -0.3264±0.0002 -0.3118±0.0015

Icorr(μA cm-2) 0.2326±0.0009 0.8833±0.0021 0.2257±0.0018 0.2146±0.0026

Ipass(μA cm-2) 4.0561±0.016 2.4615±0.019 3.1956±0.003 1.7139±0.011

ΔE 0.4353 1.1998 0.6634 0.9426

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