Effect of Microstructural Bands on the Localized Corrosion of Laser Surface-Melted 316L Stainless Steel

A recent paper¹  compares the microstructures of directed energy deposition additively manufactured (DED-AM) 316L (UNS S31603⁽¹⁾) stainless steel (SS) and of laser surface melted (LSM) 316L SS. The microstructure within each hatch of an additively manufactured 316L SS cylinder is the same as that within a single track formed by linearly translating a plate beneath a stationary laser beam. Single tracks are therefore used to investigate the influence of laser beam power and translation speed on microstructure.¹ 

It is well known that high laser power density (LPD) and high linear energy density (LED) are beneficial in completely melting the powder and creating high density deposits.^2-4  However, too high values of LPD and LED cause voids and changes in composition caused by different rates of vaporization of different elements.^4-6  In addition, high values of LPD and LED affect the microstructure of selective laser melting (SLM) deposits and additively manufactured (AM) components.¹  In the present paper, we consider high values of LPD and LED that are not high enough to create voids and significant vaporization-induced changes in composition and instead focus on the microstructures associated with high LPD and LED.

Figure 1(a) shows that the microstructures of hatches and tracks of 316L stainless steels formed with relatively high LPD and high LED are significantly inhomogeneous, which raises concerns about local variations in mechanical properties^3-4  and of other microstructure-sensitive properties such as corrosion resistance. The distinctive features of the microstructure’s inhomogeneous character are two types of bands that are revealed by metallographic etching. “Solidification bands” are so named because they form the interface between the surface melted track and the substrate, and the interface between adjacent hatches and because their mode of solidification is different from the solidification mode of the bulk of the laser melted track/hatch.¹  A second type of band, called “convection bands,” form within hatches/tracks as a result of high values of LPD and LED.¹  The convection bands appear as concentric, quasi-circular arcs in the cross section of a hatch/track (Figure 1[a]) and form half-circular ripples on the free surface of a hatch/track (Figure 1[b]).

Convection bands are the result of convection induced in the melt at high values of LPD and LED by a combination of the Marangoni effect and, especially, a high recoil pressure,⁷  and the number of convection bands increases with increasing LPD and LED.¹  The convection sweeps away the liquid at the liquid/solid interface and thereby replaces the solute-enriched interfacial liquid layer (if the solute’s equilibrium partition coefficient, k, is less than 1) or solute impoverished interfacial liquid layer (if k is greater than 1) with the composition of the bulk liquid. The abrupt change in composition of the interfacial liquid (as well as possible abrupt changes in interfacial undercooling and cooling rate) generate the abrupt changes in the mode of solidification and in the composition of the just-frozen solid, which are responsible for the differential etching that reveals the convection bands. The details of the changes in composition and in mode of solidification associated with convection bands are only apparent at high magnifications and are presented in the Results section of this paper.

It is possible that changes in mode of solidification associated with banding in LSM tracks and in AM hatches can affect the corrosion resistance. The corrosion behavior of a number of AM alloys is reviewed in a recent article.⁸  The results of a number of individual studies of corrosion of AM 316L stainless steel are described in the Discussion. In the present paper, the localized corrosion of LSM 316L SS is investigated. The results reveal the influence of banding on the susceptibility to pitting corrosion.

The surface of a plate of 316L SS was remelted by linearly translating the plate under a stationary, perpendicularly incident laser beam. The speed of translation is 5 mm/s and the power of the laser is 187 W. Figure 2 shows approximately three of a total of approximately 100, parallel, partially overlapping, laser melted tracks that cover the sample’s surface. Each track is approximately 250 μm wide and 2 cm long and the total area of melting is 2 cm × 2 cm. The laser surface melting was performed inside a glove box that contains an Aerotech XY-translation stage and a welding head (IPG Model P30-002417-V4A1A^†) of a 500W IPG YLR-500-AC-Y14^† fiber laser (1.07 μm). The atmosphere inside the glove box is stagnant ultra-high purity grade (UHP) argon. The diameter of the laser beam at the 100 mm focal length lens is 13 mm and the calculated radius (e⁻¹) of the focused Gaussian beam is approximately 10 μm. Given the surface roughness and slight curvature of the 316L SS plate, it is estimated that the laser spot on the plate’s surface is within ±1.5 mm above/below the focal spot so that the radius of the laser beam on the sample’s surface is in the range of 10 μm to 13 μm. Based on laser beam radii of 10 μm and 13 μm, and assuming the ratio of absorbed laser intensity to incident laser intensity is 0.3,^9-10  the values of the main processing parameters are LPD and LED equal 17.9 MW/cm² and 37.5 MJ/cm, and 10.6 MW/cm² and 22.2 MJ/cm, respectively.

The LSM tracks were first examined in an optical microscope and then in a scanning electron microscope (SEM). Next, the top surface and the cross-sectional surface of the laser-melted surface were ground, polished, and etched and their microstructure was examined with an optical microscope.

The etching procedure used to investigate the microstructure consisted of anodic polarization in concentrated nitric acid (68 w/o – 69 w/o, Fisher Scientific) for 10 s at a potential of +0.8 V vs. a platinum electrode, which serves as both a counter electrode and as a pseudo-reference electrode.

High-resolution energy-dispersive analysis was conducted on samples of a LSM track in an effort to determine the spatial variation of the alloy’s composition. The measurements were obtained with an FEI TitanX-60-300^† microscope. The TEM sample was prepared by focused ion beam (FIB, FEI Strata 235^†) lift out technique which is followed by nano milling (Fischione NanoMill^†) to obtain desired sample thickness.

The corrosion behavior of the plate of wrought 316L SS and of the laser surface-melted 316L SS is investigated by potentiodynamic anodic polarization in UHP nitrogen-saturated 0.1 M HCl at room temperature. The tests were conducted in a glass cell with approximately 0.5 L of solution. The samples were polarized with the aid of a potentiostat (Biologic VSP^†) at a rate of 1 mV/s. The sample’s potential is measured with respect to a saturated calomel reference electrode and is reported versus a reversible hydrogen electrode (RHE).

Samples for polarization tests were prepared by first grinding and polishing the surface. Next, the test area of the sample was created by coating the untested area with Glyptal^†, which is an electrically insulating paint. After the polarization test, the paint was removed by acetone, followed by immediate rinsing with highly purified water.

The influence of microstructure on the initiation of localized corrosion is investigated by examining in an optical microscope and a SEM (Zeiss Zemini Ultra 55^†) the surfaces of samples that are slightly etched after their local corrosion in the potentiodynamic anodic polarization tests.

The initial tests were conducted on as-laser-melted surfaces without any additional preparation. The results, which are presented in Appendix  B, indicate that topography of the as-melted surface has a dominant effect on initiation of localized corrosion and the surface topography masks the influence of microstructure. While the influence of surface topography is important, the focus of this paper is the influence of microstructure. Accordingly, the results of surface topography are presented in Appendix  B and the balance of this paper discusses the results of anodic polarization tests of metallographically polished surfaces of the plate of wrought 316L SS and of the LSM 316L SS.

Figure 3 compares the anodic polarization behavior of the ground and polished surface of the plate of wrought 316L SS and of the ground and polished LSM 316L SS surface. The polarization behavior of both the top surface and of the cross-sectional surface of the LSM layer are similar and the LSM 316L SS exhibits a significantly lower pitting potential (–0.4 V) compared to that of the wrought 316L plate (0 V). Comparison of the results presented in Figures B1 and 3 indicates grinding and polishing increases the pitting potential of the plate of the wrought 316L SS but decreases the pitting potentials of the LSM 316L SS. Although their breakdown potentials are similar, different types of localized corrosion occur on the polished top surface and cross-sectional surface of the LSM 316L SS. First, the localized corrosion of the polished top surface is described.

As is the case in most, if not all, earlier studies quasi-circular pits occur at random on the surface of the polished, wrought plate of 316L. In contrast, as shown in Figure 4(a), the pits on the polished, top surface of the LSM 316L SS have a distinctive, noncircular shape. Light etching after anodic polarization (Figure 4[b]) indicates the pits occur at specific locations in the microstructure and the shapes of the pits are influenced by the microstructure. Pits are seen to nucleate along the outer perimeter of the convection bands (Figure 4[b]). The location of pit nucleation is further revealed by the combination of the optical micrograph (Figure 5[a]) and the SEMs (Figures 5[b] and [c]) of the same region of a sample that was lightly etched after anodic polarization.

As discussed in Appendix  A, potentiostatic etching in concentrated nitric acid preferentially attacks the regions of high chromium concentration. The preferential etching is a consequence of the oxidation of Cr³⁺ to Cr⁶⁺, which is highly soluble. In Figure 5(c) the light etching indicates the initial mode of solidification along the extreme outer edge of the convection band is plane front solidification (PFS) and its smaller degree of etching indicates a lower chromium concentration, which indicates the primary phase of solidification is austenite, γ. The solidification microstructure is labeled γ-PFS. At a short distance in from the outer edge the mode of solidification changes to cellular (CS) with preferential etching of Cr-rich cell boundaries, which indicates the primary phase of solidification is austenite, γ. Here the solidification microstructure is designated γ-CS.

Ferrite is the thermodynamically stable primary phase of solidification of 316L SS and austenite is the thermodynamically metastable primary phase.^11-14  The initial freezing as PFS of the metastable austentite phase (γ-PFS) followed by austenitic cellular solidification (γ-CS) occurs as a result of the high cooling rate and high undercooling at the outer edge of the convection band.¹  The change from plane front solidification to cellular solidification is the result of transitioning from transient to steady-state conditions¹⁵  and/or a consequence of the decrease in G/V (= thermal gradient in the liquid at the liquid/solid interface divided by the velocity of the liquid/solid interface) with increasing distance from the fusion line.^16-18  The SEM in Figure 5(c) of the lightly etched, pitted sample shows that submicrometer diameter pits (approximately 0.2 μm) nucleate within the cores of the austenite cells. In addition, the SEMs in Figures 6(a) and (b) of the lightly etched, pitted sample show submicrometer diameter pits in both the γ-PFS region as well as in the austenite cores of the γ-CS region. The diameter of the solidification cells is approximately 1 μm, which is clearly larger than the 0.2 μm diameter of the tiny pits that appear homogeneously distributed throughout the interior of the cell (Figures 5[c] and 6[b]). Additional TEM work, which is outside the scope of the current investigation, is required in order to determine if there is a microstructural cause of the 0.2 μm scale of the pits.

Significant changes in the microstructure occur with increasing distance from the two narrow regions of primary austenite solidification (i.e., γ-PFS and γ-CS) that occur along the outer edge of the convection band. The changes in microstructure consist of changes in the primary phase of solidification and changes in the mode of solidification. However, pits do not initiate in any of these regions. Because pitting corrosion is the primary focus of this paper the additional changes in microstructure are presented and discussed in an appendix (Appendix  C).

The entire microstructure is composed of a series of convection bands and the microstructure within each band is summarized in the sketch presented in Figure 7. Starting at the outer edge of the convection band the initial solidification occurs as γ-PFS followed by γ-CS. Next, as discussed in Appendix  C, there is a transition from γ-CS to α-CS followed by a transition from α-CS to α-equiaxed dendritic solidification. Although the regions of primary austenite solidification are very small compared to the relatively large regions of primary ferrite solidification, the regions of primary austenite solidification are very important as they are the sites where pits are nucleated.

As shown in the SEM of Figure 8(a), localized corrosion of the cross-sectional surface initiated at the interface between the paint mask and the sample. The sizes of the sites of localized corrosion are very different on the top surface and cross-sectional surface. As already discussed in the Summary of Microstructure and Localized Corrosion section, on the top surface a large number of relatively small pits form along the outer edges of convection bands. In contrast, the localized corrosion of the cross-sectional surface consists of a few, relatively large regions of crevice corrosion.

The morphology of the interior surface of the crevice corrosion shown in Figure 8 is revealed by the higher magnification micrograph presented in Figure 8(b). At the outer edge of the crevice corroded region (lefthand side of Figure 8[b]) the interiors of the solidification cells are more strongly attacked leaving the cell boundaries raised in relief. The surface morphology suggests the mode of solidification along the perimeter of the crevice corroded region is γ-CS. We learned from metallography of the top surface that γ-CS is the mode of solidification along the outer edge of the convection band. Thus, the perimeter of the site of crevice corrosion pictured in Figure 8 occurs along the outer edge of the convection band, but we believe this is coincidental.

For clarification in our view there is no specific reason why crevice corrosion would stop at the edge of a convection band as it has in Figures 8(a) and (b). In fact, the volume of each site of crevice corrosion is relatively large, indicating that the crevice corrosion has propagated through a number of convection bands. The main conclusion of the surface morphology is that within the occluded volume of the crevice the rate of corrosion of the chromium-depleted austenitic cores is greater than the rate of corrosion of the chromium-enriched cell boundaries.

As summarized in the sketch presented in Figure 7, further in from the edge of the convection band the primary phase of solidification changes to ferrite. As indicated in Appendix  C1, the transition from primary austenite to primary ferrite occurs at a distance of approximately 10 μm from the outer edge of the convection band. At approximately the same distance of 10 μm from the edge of the crevice corrosion shown in Figure 8(b) the morphology of the crevice corroded surface changes. The change in surface morphology corresponds to the transition from γ-CS to α-CS.

The morphology of the corroded surface of α-CS, which is presented in the righthand side of Figure 8(b), suggests that unlike the case of γ-CS, there is not a significant difference in the amount of corrosion of the primary ferrite cell’s interior and boundary. The morphology of the crevice corroded surface along with the previously mentioned absence of pit initiation in the regions of α-CS suggest that the lower chromium concentration of the cell boundary is high enough to provide corrosion resistance equivalent to that of the cell’s core. The higher chromium concentration of the primary ferrite cells is a consequence of the fact that chromium’s equilibrium partition coefficient is greater than one for the ferrite phase.^11-12  It is also possible that the distribution of chromium within the primary ferrite cell is homogenized as a consequence of the high solute diffusivity of the body-centered cubic crystal structure.

Our results indicate that the mode of solidification and the primary phase of solidification, which are sketched in Figure 9 along with information about the spatial extent of each zone of solidification, strongly affect the initiation of localized corrosion. Specifically, on the top surface of ground and polished, LSM 316L SS pits preferentially initiate along the outer edge of the convection band, which is the region of transient solidification, and which is composed of the low chromium, primary austenite zones of plane front solidification and cellular solidification. Because molybdenum is also a ferrite stabilizer like chromium it is possible that the austenitic region is also depleted in molybdenum and that the combination of chromium and molybdenum depletion is responsible for the nucleation of pits in the primary austenitic regions along the outer edge of the convection bands.

The sites of crevice corrosion, which occurs during anodic polarization of the cross-sectional surface, are much larger than the pits formed during anodic polarization of the top surface and the volume of each site of crevice corrosion is composed of multiple convection bands. The surface morphology of one site of crevice corrosion reveals the primary austenitic mode of solidification along the bottom of the laser remelted volume and the primary ferritic mode of solidification within the interior of the laser remelted volume.

It is noteworthy that localized corrosion is not initiated in the region of primary ferrite solidification. Thus, our results indicate the steady-state solidification structure (i.e., primary ferritic, cellular solidification) is more resistant to pitting corrosion than is the narrow zone of the initial, transient solidification of γ-PFS and γ-CS. That is, the increased susceptibility to pitting corrosion of LSM 316L SS is attributed to the transient stage of solidification in which the first phase to freeze is the low-chromium-containing metastable austenite.

It is instructive to compare the present results with the results of several earlier studies of the pitting resistance of wrought and rapidly solidified (AM and LSM) 316L. Previous studies indicate two aspects of microstructure of 300 grade stainless steels that affect resistance to pitting corrosion: (1) inclusions and (2) phase transformations that cause local changes in concentrations of chromium and molybdenum.

First, consider inclusions. MnS inclusions are known to preferentially initiate pits in wrought 300 grade stainless steels.^19,20  In some studies in which AM exhibits greater resistance to pitting than the wrought material, the superior resistance of AM is correlated with the absence or reduced size of MnS inclusions. For example, Stewart and Williams²¹  reported that LSM increased the pitting resistance of 304L in dilute neutral chloride by markedly decreasing the size of sulfur-containing inclusions. Sander, et al.,²²  found that the pitting potential of laser powder bed fusion (AM-LPBF) 316L is greater than that of wrought 316L in 0.1 M NaCl and they suggested the greater pit resistance of AM is a consequence of the absence of, or reduced size of MnS inclusions. In a closely similar conclusion, Chao, et al.,²³  determined that the higher pitting potential of AM-LPBF 316L in 0.6 M NaCl (approximately 0.5 V higher) compared to that of wrought 316L is a consequence of the complete absence of MnS inclusions in the AM microstructure.

In contrast to the results of Stewart and Williams and of Sander, et al., Ganesh, et al.,²⁴  and Trelewicz, et al.,²⁵  indicate the pitting resistances of AM-DED and AM-LPBF, respectively, are less than that of wrought 316L in 3.5% NaCl and in 0.1 M HCl, respectively. In both cases the lower pit resistance of AM is attributed to phase transformations that create localized regions of low chromium and low molybdenum concentrations. Relatively low-magnification optical microscopy in Ganesh, et al., indicates that in relatively large diameter pits (i.e., 20 µm to 30 μm) the cores of primary austenite cells are preferentially corroded, and the cell boundaries are uncorroded. In addition, the γ−χ grain boundaries are also preferentially pitted due to sensitization. Solution annealing (SA) healed the chromium-depleted zones of the sensitized grain boundaries so that pits did not form along γ−γ grain boundaries. SA did not homogenize the Cr concentration of the primary austenite cells so that within the pits the cell cores were preferentially corroded, and the cell boundaries were uncorroded as is the case for the as-LSM samples.

In instances of primary ferritic AM Ganesh, et al., found that pits preferentially formed along austenite-ferrite boundaries, which they suggested is due to segregation of sulfur and phosphorus, and along sensitized γ−γ grain boundaries. As was the case for primary austenite AM, SA removed the Cr-depletion along γ−γ grain boundaries so that pitting did not occur along γ−γ grain boundaries of SA primary ferritic. Trelewicz, et al., does not include a microscopic examination of the surfaces of pitted samples. But their microchemical analyses indicate that molybdenum preferentially segregates to the cell boundaries of primary austenite and they suggest that the lower pitting resistance of AM is a consequence of low molybdenum concentration of the cores of primary austenite cells.

Finally, it is worth mentioning several additional comparative studies if only because they provide futher evidence of the wide variation in the relative resistances to pitting corrosion of wrought and rapidly solidified (i.e., AM and LSM) 316L. Lohdi, et al.,²⁶  found that AM exhibits greater pitting resistance than wrought 316L in 0.9% NaCl, but in two other studies conducted in the same electrolyte Sun, et al.,²⁷  and Zietala, et al.,²⁸  indicated that AM-DED 316L has lower pitting resistance than wrought 316L on account of porosity and because of chromium segregation, respectively. Majumdar, et al.,²⁹  measured the pitting potential of AM-DED cladding of 316L in 3.56% NaCl to be 125 mV higher than that of wrought 316L. Lohdi, et al.,²⁶  and Majumdar, et al.,²⁹  did not identify the initiation sites of pitting and the morphology of corroded pits. Lohdi, et al.,³⁰  found that neither AM nor wrought 316L pit during potentiodynamic polarization in acidic solutions (pH 1 to 3) with dilute concentration of chloride (≈0.01 M).

Thus, our results, which indicate the pit resistance of LSM 316L is less than that of wrought 316L and indicate the lower pitting potentials of AM is due to the formation of very narrow regions of primary austenite, are consistent with the results of Ganesh, et al., and Trelewicz, et al., which also found or suggested corrosion preferentially occurs in low Cr regions, i.e., cores of primary austenite as well as along sensitized γ−γ grain boundaries of AM.

Therefore, our results in combination with earlier investigations of the comparative pitting corrosion of wrought and rapidly solidified (AM and LSM) 316L indicate that in those cases in which rapidly solidified 316L has greater resistance to pitting than wrought 316L, the greater resistance is attributed to the absence of or reduced size of MnS. And, in those cases in which rapid solidification exhibits less resistance to pitting than wrought, the lower resistance is due to the presence in the rapidly solidified microstructure of chromium-depleted zones, e.g., the chromium depletion associated with sensitization and with the low chromium concentrations of cores of primary austenite cells.

• On ground and polished surfaces of laser surface melted 316L SS pitting corrosion preferentially initiates at convection bands that are created by a combination of the Marangoni effect and a high recoil pressure.

• Within the convection bands the mode of solidification changes with increasing distance from the outer edge. γ-plane front solidification (PFS) occurs along the extreme outer edge and changes to γ-cellular solidification (CS), and then to two regions of primary ferrite solidification: α-CS followed by α-equiaxed dendritic solidification (EDS).

• Pits initiate in the regions of primary austenite solidification along the outer edge of the convection band. Specifically, the nucleation of pitting corrosion occurs in the very narrow region of transient solidification along the edges of the convection bands. The steady-state microstructure of 316L SS consists of α-CS and α-EDS and is more resistant to pit initiation than is the transient γ-PFS and γ-CS.

• Sites of crevice corrosion created during anodic polarization of the cross-sectional surface grow through multiple convection bands. The morphology of the surface of crevice corrosion indicates preferential corrosion of the chromium-depleted austenitic cores of the solidification cells, which places the chromium-rich intercellular regions in positive relief. The morphology of the surfaces of crevice corrosion of the region of α-CS solidification indicate there is no significant difference in corrosion attack of the cell’s cores and the cell boundaries of α−CS, which is the steady-state mode of solidification of 316L SS.

• Anodic polarization in aqueous 0.1 M HCl of unpolished, as-laser-melted surfaces indicates that the high laser power used (≈10 MW/cm² to 18 MW/cm² absorbed power density and ≈ 22 MJ/cm to 38 MJ/cm linear energy density) creates metal droplets on the free surface that nucleate pitting corrosion.

• The combination of the present results and the results of earlier studies suggests that instances in which rapidly solidified 316L exhibits greater susceptibility to pitting corrosion than wrought 316L are a consequence of phase transformations (sensitization and solidification) that create local regions of chromium-depleted zones in the rapidly solidified microstructure. And, in cases in which rapidly solidified 316L exhibits greater resistance to pitting corrosion than wrought 316L, the rapidly solidified microstructure is free of local regions of chromium depletion and is either free of MnS inclusions or contains MnS inclusions that are very small and below the critical size required to initiate stable corrosion pits.

We are grateful to the Molecular Foundry and National Center of Electron Microscopy at the Lawrence Berkeley National Laboratory for the use of the Zeiss Gemini Ultra-55 analytical scanning electron microscope, the Bruker AXS D8 Discover GADDS x-ray diffractometer, the FIB Strata, the Nano Mill and the FEI TitanX 60-300. Work at the Molecular Foundry was supported by the Office of Science, Office of Basic Energy Sciences, of the U.S. Department of Energy under contract no. DE-AC02-05CH11231. Our research was made possible by the programmatic support of Dr. Andy Vance. Sandia National Laboratories is a multimission laboratory managed and operated by National Technology and Engineering Solutions of Sandia, LLC., a wholly owned subsidiary of Honeywell International, Inc., for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-NA0003525. This paper describes objective technical results and analysis. Any subjective views or opinions that might be expressed in the paper do not necessarily represent the views of the U.S. Department of Energy or the United States Government.

In this investigation etching is conducted potentiostatically in concentrated nitric acid at an applied potential of +0.8 V_RHE. The following results demonstrate that our etching procedure preferentially attacks chromium-rich regions. Anodic polarization curves were measured of three binary iron-chromium alloys: Fe-5w/oCr, Fe-10w/o Cr, and Fe-13 w/o Cr as well as of Fe-5 w/o Cr – 10 w/o. Each of the four alloys is cast in vacuum as a 19 lb cylinder, followed by hot and cold working into 1.52 mm diameter wires. Each sample of the anodic polarization test is approximately 40 mm in length. Each test is conducted in approximately 500 cc of air-saturated, concentrated nitric acid (65 w/o). Platinum mesh is the counter electrode, and a 2 mm diameter platinum wire serves as a pseudo-reference electrode. The potential of the platinum pseudo-reference electrode in air-saturated concentrated nitric acid is +1.108 V with respect to the saturated calomel electrode. Each sample is anodically polarized at a rate of 1 mV/s.

The anodic polarization curves of the four alloys are presented in Figure A1. The sharp increase of oxidation rate in the anodic polarization curves at potentials of approximately +0.45 V, +0.60 V, and +0.65 V with respect to the platinum pseudo-reference for Fe-13Cr, Fe-10Cr, and Fe-5Cr, respectively, is the result of oxidative dissolution of Cr₂O₃ to hexavalent chromium. The polarization behavior of Fe-5Cr-10Ni indicates that 10 w/o Ni does not have a significant effect on the potential at which Cr₂O₃ is oxidatively dissolved.

The results presented in Figure A2(a) through (d) confirm that our etching practice preferentially attacks chromium-rich regions. In the cellularly solidified portion of the laser remelt track of 316L SS (Figure A2[a]), which is formed with a power density of 200W and a translation speed of 4.3 mm/s, the etching procedure preferentially attacks the intercellular regions, which the elemental maps of nickel, molybdenum, and chromium (Figures A2[b] through [d]) and the quantitative analyses of the intercellular region (see Table A1) indicate are chromium-rich.

The anodic polarization curves of the as-received plate of 316L SS and of the as laser-remelted surface are presented in Figure B1.

The polarization curves of both the as-received plate and the as-laser-remelted surface are similar. However, whereas the pits were randomly distributed on the surface of the as-received plate, the pits were preferentially located on the portion of each track that was immediately adjacent to the next-deposited track. As shown in Figure B2(a) and explained by the sketch in Figure B2(b) the pits occurred in the region of the track that was partially coated with molten-alloy droplets created by the recoil pressure.⁷ 

EDS spectra indicated there are no significant differences in composition in the different regions of the track. A possible cause of the preferential pit location is initiation by underdeposit corrosion at the site of the metal droplets. Comparing the results of the as-received plate and the laser-remelted surface indicate that although there is no difference in the pitting potentials, there is a significant increase in the number of pits associated with the as-laser-remelted surface.

From the perspective of initiation of pitting corrosion, the most important microstructure is the two regions of primary austenite solidification (i.e., γ-PFS and γ-CS) that occur along the outer edge of the convection bands. A series of changes in microstructure occur beyond this region of primary γ solidification. We describe these changes for the sake of completeness, but the results and discussion are presented in this appendix rather than in the body of the paper because these changes do not affect the initiation of localized corrosion. It is possible that these changes have an impact on other types of corrosion, such as environmentally assisted cracking.

Information about the influence of microstructure on the initiation of pitting corrosion is provided by searching for evidence of corrosion at increased distance from the edge of the convection band, which is where initiation occurs. Figure C1(a) through (c) present optical micrographs of a more heavily etched surface than that shown in Figures 7 and 8. The relatively low magnification of the micrographs in Figures C1(a) and (b) reveal the cyclic changes in microstructure associated with the convection bands. Note that each convection band is composed of two sub-bands, e.g., A and B, and C and D. One sub-band (e.g., A and C) is darker as a consequence of more etchable features. The other sub-band has few etchable features and is light colored (e.g., B and D). The microstructure along the outer edge of the sub-bands A and C is not resolvable given the low magnification of the optical micrographs but consists of the narrow regions of γ-PFS and γ-CS shown in Figure 7(c).

The vast majority of sub-bands A and C is composed of α-CS as is apparent in the magnified image in Figure C1(c). The cell boundaries and the cells’ cores are identified in Figure C1(c). The solidification microstructure is identified as α-CS on account of (1) the continuous, lightly etched regions that are in relief and (2) the deeper attack of the island-like regions. The continuous regions form the cell boundaries, and the island-like regions form the cores of the cells. The deeper attack of the cores indicates the cores have a relatively high chromium concentration and, therefore, solidified as ferrite. The lighter attack of the cell boundaries indicates a relatively low chromium concentration.

The change in the primary phase of solidification from austenite to ferrite at a distance of 10 μm from the outer edge of the convection band (i.e., the outer edge of sub-band A) is a consequence of the following: (i) once quasi steady-state solidification is reached the undercooling associated with a high laser power such as 200 W is relatively small and (ii) low undercooling promotes freezing of the thermodynamically stable phase, which is ferrite for 316L SS.

The boundary between the dark sub-band and the light sub-band occurs at a distance of approximately 50 μm from the edge of the convection band and marks a change in the mode of solidification from α-CS to α-ED (equiaxed dendritic) (see Appendix  C2).

It is worth noting that the light sub bands are not sites of initiation of pitting. Although there is no effect of the light sub bands on pitting, which is the primary focus of this paper, the different microstructure of the light sub band might affect resistance to EAC and so this microstructure is further discussed in Appendix  C2.

The influence of microstructure on the localized corrosion is summarized in the sketch presented in Figure 9 in the body of the paper.

The cross section of the laser-surface melted track formed with a laser power of 375W and a translation speed of 25 mm/s is presented in Figure C2. A narrow band of plane front solidification is present at the interface of the deposit and the substrate. Approximately 5 µm from the interface with the substrate the mode of solidification changes to cellular. The chromium-rich regions are preferentially etched and form the cells’ boundaries, which indicates austenite (face-centered cubic) is the primary phase of solidification. As distance from the substrate increases the mode of solidification changes from cellular to equiaxed-dendritic. The region of equiaxed grains is associated with high laser power, which creates a high recoil pressure. The high recoil pressure induces significant convection of the melt pool. The high convection breaks off pieces of cells that then serve as nuclei of solidification. The transition from cellular solidification to equiaxed dendritic solidification occurs along a very rough, irregularly shaped boundary.

Refer once again to the micrograph that is presented in Figure B1(b). Note the irregular shape of the sharp interface between the regions of α-CS and α-EDS. The EDS zone labeled “D” is wider than the others that are shown and contains two layers of grains. It is also worth noting that the nitric acid etch indicates there is less variation in chromium concentration between the cores and the interdendritic spaces. In other words, the chromium concentration is more uniform, which might account for the absence of initiation of pitting corrosion in the α-EDS regions.