Selective laser melting (SLM) or powder bed fusion is a type of additive manufacturing technology with applications in, e.g., the orthopedics, energy, and aerospace industries. Several studies investigated the localized corrosion behavior of SLM-fabricated Type 316L (UNS S31603) stainless steel. However, little is known about the effects of tribocorrosive conditions on the response of stainless steels fabricated by SLM. In this study, the effects of third-body abrasive particles on the tribo-electrochemical behavior of SLM 316L stainless steel produced by SLM were investigated and compared with wrought counterparts (including UNS S31703, 317W) in 0.6 M NaCl. It was found that the presence of Mo played a more decisive role in the tribocorrosion behavior than the manufacturing method, i.e., 317W revealed the best tribocorrosion behavior vis-a-vis wrought 316L and the SLM-fabricated specimens. The improved tribocorrosion behavior contrasted with the much higher breakdown potential of the SLM-fabricated samples. Nano-scale secondary ion mass spectroscopy was used to investigate the effects of Mo on passivity. The implications of passivity and tribocorrosion behavior are discussed.
Stainless steels’ excellent corrosion resistance relies on the formation of a nanoscale surface passive film,1 which drastically slows down the uniform dissolution of the underlying substrate.2 The nanoscale passive film is, however, vulnerable to chemical and mechanical breakdown.3 Passivity breakdown occurs due to local acidification within a small cavity or crevice, leading to stable localized dissolution4 and possible catastrophic failures. Hence, localized corrosion is of great concern in different industries.5
Several studies aiming at predicting the mechanism and prevalent conditions leading to stable localized dissolution—e.g., pitting, crevice, or intergranular corrosion—are available in the literature.6-8 It has been shown that chloride ions, for example, decrease the induction time for pitting or crevice corrosion to occur.9 Likewise, detrimental effects of impurities and inclusions, e.g., MnS inclusions in the case of wrought 316L (UNS S31603(1)) stainless steel, on pitting corrosion have been discussed.10-12 In addition to chemical means of film breakdown, surface contact with hard grit particles can disrupt passive films under abrasive and erosive conditions.13-14 In particular, the tangential movement of loaded particles against the surface can damage the passive film, plastically deform the surface and lead to direct material removal.15 Fast repassivation after passivity breakdown is, thus, paramount to minimize stable pit and crevice propagation. Therefore, many studies have been conducted to determine the critical conditions for stainless steel repassivation.16
The different manufacturing routes and the resulting microstructure appear to have a crucial role in the electrochemical-mechanical behavior of stainless steels.17-18 Selective laser melting (SLM)—or powder-bed laser fusion additive manufacturing—allows layer-by-layer fabrication of solid metallic parts and components. Type 316L stainless steel is among the alloys that can be manufactured with SLM, and it has been the subject of many investigations.19-24 As the SLM process involves localized metal melting, the microstructure achieved by SLM can be very complex; hence, predicting their corrosion response can be challenging.
In the past few years, several studies have been performed to determine the microstructure and corrosion behavior of SLM-manufactured 316L (SLM 316L) stainless steels as some reviewed in Schindelholz, et al.18,25 Wang, et al.,26 showed the segregation of Mo at the boundaries of cellular networks formed within the alloy’s microstructure. Effects of heat treatment on mechanical properties and microstructural evolution of SLM-fabricated stainless steels have been also studied.19,27-28 Recently Chao, et al.,29 investigated the effects of heat treatment on the microstructural evolution of SS316L additively manufactured via SLM. Their study indicated that annealing at 1,000°C resulted in the recrystallization, grain growth, and coarsening of nano inclusions.
Regarding the corrosion behavior of additive manufacturing (AM) stainless steels, some studies claimed superior corrosion resistance vis-à-vis the wrought counterpart.21-22,27,30-32 In this regard, the critical pitting temperature (CPT) of wrought SS316L was found to be 16°C, whereas pitting corrosion of the SLM-fabricated alloy was observed above 27°C to 31°C.33 It was suggested that SLM-manufactured SS316L had a wider passive region owing to microstructural grain refinement. Researchers suggested that AM improved the passivity of the material due to the formation of a nanoscale cellular structure.34-35 On the other hand, unexpected decline in pitting resistance following heat treatment36 and unusual intergranular corrosion behavior of AM SS316L has also been reported without37 and with postprocessing.38 Mo and Cr segregation at the subgrain boundaries has been discussed.32 Other studies attributed the higher dissolution at the melt pool boundaries to residual stresses and microstructural inhomogeneity.39
On the effects of a mechanical contact on depassivation of these alloys, a study has been performed on the erosion-corrosion behavior as a result of impinging particles.40 In addition, specimens under a two-body sliding tribocontact were investigated.41 However, possible detrimental effects of mechanical abrasion on passive film removal are yet to be studied.
In this work, the effects of third-body silica sand particles on the three-body abrasion-corrosion behavior of SLM 316L stainless steel were studied under different heat treatments and compared with wrought UNS S31603 and S31703. This study focused on determining whether stable localized corrosion would occur following mechanical depassivation by hard abrasive particles. The two wrought materials tested had a different Mo content, which is an important element in hindering pitting corrosion.42 Herein, the effects of Mo on the tribocorrosion behavior of austenitic stainless steels are discussed for the first time. Tribo-electrochemical methods combined with characterization techniques were used to determine their repassivation behavior post abrasion at different anodic potentials.
Type 316L stainless steel samples were additively manufactured using a 3D systems Pro X† DMP 320 at a layer thickness of 30 µm with a scanning speed of 900 mm/s and a laser power of 250 W. The 316L powder used for the SLM fabrication was supplied by TLS Technik GmbH & Co., Germany. The chemical composition of the powder was shown in Table 1. The powder size was 10 µm to 45 µm. Following the SLM fabrication process, a group of samples was solution annealed at 1,050°C for 4 h in an argon-purged tube furnace. Subsequently, the samples were either water quenched, or furnace cooled.
Commercially sourced (Outokumpu) wrought Type 316L and 317L (UNS S31603 and S31703) austenitic stainless steels were annealed at 1,050°C for 4 h and furnace cooled. Their chemical composition is also given in Table 1.
In this study, the samples were named: 316AP: as printed Type 316L (i.e., as SLM manufactured), 316AM-HT-FC: SLM additively manufactured Type 316L, followed by annealing and furnace cooling, 316AM-HT-WQ: additively manufactured, then annealed and water quenched, 316W: UNS S31603, annealed, and 317W: wrought UNS S31703, annealed.
The microstructure of the samples was examined after electrolytic etching with 10 wt% oxalic acid, using optical and scanning electron microscopy (SEM). Prior to etching, the samples were polished down to a mirror finish (Ra = 0.01 µm) using a 1 µm diamond suspension. An FEI Talos† FS200X G2 transmission electron microscope (TEM) equipped with energy-dispersive x-ray spectroscopy (EDS) was used to characterize the microstructural features of the SLM-manufactured samples before and after heat treatments. For the EDS measurements, a 25 kV beam voltage was used. Site-specific TEM lamellae were made by Tescan† Lyra3 focused ion beam (FIB) coupled with SEM. In addition, FIB cross sectioning was used to analyze the areas of interest. Prior to FIB milling, a platinum strip was first deposited over the region of interest. Finally, 1 nA ion beam current was used to smooth the newly created cross-section surface. The surface was finally examined with both secondary electron (SE) and ion beam channeling contrast (ICC) imaging. The latter was used to analyze grain miss-orientation below the worn surfaces. Tescan Clara and Mira field emission SEM was also used to image the samples’ surface after abrasion-corrosion tests.
CAMECA† NanoSIMS 50 and NanoSIMS 50L were used to characterize the passive film formed on two types of 316W with 2% and 2.5% Mo in the microstructure and 317W sample with 3.1 wt% Mo and also the additively manufactured 316L. The passive film was formed by exposing the ground specimen to air for at least 24 h. The ionic distributions were measured as a function of passive film depth. A modified depth-profiling procedure as described by Salasi, et al.,43 was used to sputter and analyze the passive film. In this method, a thin Cs layer was deposited on the sample surface inside the NanoSIMS chamber by operating the Cs primary beam at low energy. This procedure preserves the passive layer and provides a sufficiently high dose of the primary ions to achieve steady-state secondary ion generation before sputtering. A Cs+ primary ion beam was used to profile the presence of 16O−, 52Cr16O−, 56Fe16O−, and 96Mo16O− species at different surface sublayers. Sputtering was repeated in triplicate on different locations of the same sample to ensure the reproducibility of the measurements, and representative results are shown in this study.
Test Rig and Conditions
An SP300† Biologic potentiostat was used for the electrochemical tests, performed in a three-electrode configuration. An Ag/AgCl (sat.) (SSE) and platinum wire were used as the reference and counter electrodes, respectively.
The corrosion-only cyclic potentiodynamic polarization tests (CPP) were performed in an EG&G flat cell in the 0.6 M NaCl. CPP tests during three-body abrasion-corrosion were conducted in the tribo-electrochemical test rig explained above (see Figure 1). The CPP tests were performed at 1 mV/s scan rate starting at −0.1 VSSE below the open-circuit potential (OCP) until reaching 5 mA or +1.0 VSSE; then, the scan was reversed.
Interrupted potentiostatic polarization tests were performed at different anodic potentials to investigate the repassivation behavior of the tested alloys following abrasion. The applied potential was held constant at the desired value for 300 s without abrasion, 60 s with abrasion, followed by 300 s of a no-abrasion rest period. The polarization started from the OCP and increased at 0.1 V steps. Once the maximum desired potential was reached, the potential was reverted to −0.1 VSSE, after which the test was terminated.
The repassivation behaviors of the wrought and SLM-manufactured 316L and wrought 317L stainless steels were also investigated using the potential pulse method (PPM).47 First, the sample’s potential was held at −1.2 VSSE for 120 s; then, the potential was increased to an anodic potential of +0.4 VSSE. Ramping up took 2 µs and the resulting current was measured simultaneously as a function of time.
Characterizing the Materials
Additionally, melt pool boundaries and a mixture of equiaxed and columnar structures were seen throughout the microstructure. Figure 2(c) also shows an austenite grain intersected by a melt pool boundary with spherical pockmarked surface dissolutions, as indicated by the arrow. As shown in Figures 2(d) through (g), annealing (irrespective of the cooling rate) led to the disappearance of melt pool boundaries, although some faint rings remained, delineating the prior melt pool lines, as indicated by the arrows. Additionally, the emergence of a more uniform microstructure—where there was no columnar or equiaxed microstructure—was apparent after annealing. Another specific feature visible in etched heat-treated samples was the formation of large irregularly shaped inclusions.
Figure 6 also shows the tribo-electrochemical data recorded while the abrasive particles were delivered and loaded against the samples. There was an abrupt current increase above anodic potentials for the Type 316L samples, which increased, in order, from 316W: +0.32 VSSE > 316AP: +0.28 VSSE > 316AM-HT-WQ: +0.21 VSSE > 316HT-FC: +0.04 VSSE, while for the 316W specimens, the current continually increased above +0.28 VSSE, the current of the 316AP sample increase was limited and secondary passivity was later established. The 316AP-WQ and 316AP-FC, however, showed no passive behavior above the potentials mentioned previously. In the case of 317W under abrasion, the sample showed pseudo-passivity along the polarized potentials up to the anodic potential of +1.0 VSSE.
Table 2 summarizes the current extracted from the polarization curves under three-body abrasion-corrosion at corrosion potential (Ecorr). Also in this Table current values at −0.2 VSSE applied potential, denoted as Ipp or a pseudo-passive current are reported. It can be seen that the printed 316L material demonstrated a lower current under abrasion under Ecorr and once polarized. The values are comparable to that of the wrought 317L (317W) specimen.
Under three-body tribocorrosion, the 316W’s anodic current measured above an applied potential of +0.28 VSSE was driven by stable pitting corrosion rather than mechanical depassivation by the silica sand particles. Interrupted tribocorrosion tests also showed that abrasion, in fact, reduced the current measured, perhaps by hindering the propagation of the active pits. Contrastingly, the SLM-fabricated samples under abrasion and anodic potentials showed active dissolution. Only the 317W with a higher Mo content remained passive along the entire polarization range. Regarding Ecorr changes under abrasion, 317W showed the highest potential drop to −0.50 (±0.01) VSSE, whereas the potential of all 316L samples shifted to −0.40 (±0.03) VSSE.
This study showed that the 316AP specimens, regardless of their heat-treatment state, had a higher Eb than that of 316W. The apparently improved localized corrosion resistance was in agreement with Chao, et al.21 SEM and TEM analyses of the microstructure confirmed the absence of MnS inclusion in the SLM 316L samples. Instead, nano-scale Mn- and Si-rich oxides were finely dispersed in their microstructure. it is known that MnS inclusions act as preferred pit nucleation sites in 316L stainless steel,10 their absence in the as SLM-fabricated samples led to a higher breakdown potential. In this study, it was found that annealing after SLM, however, caused agglomeration of the inclusions and the formation of secondary phases followed by a decrease in the Eb. Laleh, et al., discussed the changes in the morphology and chemistry of nano-scale inclusions following heat treatment.36 They found that heat treatment resulted in MnS-rich inclusions that adversely affected localized corrosion resistance. Herein, we found no MnS-rich inclusions postheat treatment, however, large irregular inclusions enriched in Cr, O, and Si were found. A plausible explanation for the postheat treatment reduced localized corrosion resistance is that irregularly shaped larger size postheat treatment inclusions could act as microcrevice sites, facilitating localized corrosion,51 although further research is needed on this. A similar effect has been reported on the effects of porosity as reviewed in Schindelholz, et al.18 Independently of the role of inclusions, the repassivation potential of the samples—which depends on the extent of pitting/crevice corrosion but it is independent of the presence of inclusions52 —were almost identical, suggesting a similar repassivation behavior regardless of the fabrication method and heat-treatment route. Likewise, PPM data confirmed a similar repassivation behavior. In other words, our findings suggested that once a pit initiates it will continue to propagate, presenting a repassivation behavior that was independent of the fabrication method.
The corrosion-only behavior contrasted with the tribocorrosion results, which indicated a clear difference between the SLM-fabricated 316L and wrought 317L stainless steel samples. Tribocorrosion tests showed that each material had a different character following a mechanical disruption of the passive film. From the tribocorrosion tests, three important observations were made:
Inclusions vs. passivity. Tribo-electrochemical tests indicated that the passive film controlled the anodic tribocorrosion behavior of the SLM-fabricated samples. On the other hand, localized corrosion of the wrought stainless steels was a substrate-driven mechanism, where MnS inclusions played a crucial role. In the heat-treated SLM 316L samples, once the sand particles disturbed the surface, the material showed active dissolution, and the current increased at relatively low anodic potentials. The current increase was not influenced by the higher Eb of the additively manufactured samples.
Effect of Mo. In this study, it was shown that the Mo content had a more marked effect in the resulting tribocorrosion resistance than the manufacturing method; in particular, at potentials over +0.30 VSSE. Previous studies indicated that Mo increased the resistance of austenitic stainless steels to pitting and crevice corrosion, specifically in chloride-containing environments.17 The hypotheses around plausible Mo effects can be summarized into two categories: (a) those discussing the effects of Mo on passivity53-54 and (b) those concentrating on pit chemistry and salt film formation.42,55 On the former (Mo and passivity), the x-ray photo electron spectroscopy (XPS) surface analysis studies revealed: (i) Mo(VI) enrichment occurs at the outermost layer of the passive film56 and (ii) the presence of Mo ions within the passive film’s barrier layer inhibits the rate of dissolution of iron and chromium from the matrix.57 In this regard, the NanoSIMS results presented herein showed a Mo-O enrichment at the inner barrier layer, regardless of the amount of Mo in the matrix, which was between 2% and 3%.
A recent study by Henderson, et al.,58 found Mo enrichment in a passive film during transpassive dissolution. An early study by Sugimoto and Sawada59 using XPS found that the passive films formed on molybdenum-containing stainless steels were composed of a complex oxyhydroxide-containing Cr(III), Fe(III), Ni(II), Mo(VI), and Cl−, which increased in thickness with increasing Mo content. They proposed that the solid solution formed between Mo(VI)- and CrOOH-inhibited Mo from dissolving transpassively, a process that usually happens at +0.27 VSSE. This indicates that there is a possible mechanism by which passivity in stainless steels can be maintained at potentials higher than the transpassive potential of Mo.58 These findings could explain the fact that even at highly anodic potentials under abrasion the passive film on the 317W steel remained stable. In the CoCrMo alloys commonly used in hip arthroplasty, it has been found that the release of Mo(VI) and the formation of complexes with proteins could reduce the tribocorrosion resistance of the material.60
Regarding Mo effects on the local pit chemistry, the SEM images and FIB-SEM analysis revealed several pits in the worn areas of the 316W but only a couple in 317W. Although pits were observed in the 317W, the current measured remained unaffected, contrasting with the 316W results. This difference could be explained by the effects of additional Mo in 317W, which could inhibit stable pit growth under the abraded surface. Studies have shown that molybdate ions within the pit developed stable complex salts with chlorides, thus, reducing the chloride concentration within the pit.61 In addition, as a result of the salt precipitated at the bottom of the pit, repassivation subsequently occurred.57
Tribo-metallurgical layer. Signs of surface delamination found on the surface after tribocorrosion tests on SLM samples could indicate the formation of a tribo-metallurgical layer, which may play an important role in their repassivation behavior. The layer could be a mixture of highly strained bulk material and inclusions. Formation of such tribolayer and alteration of near-surface structure during tribocorrosion has been reported before.14-15,62-63 Further research is needed to investigate the implications of this layer on the repassivation behavior of the material.
Based on the comprehensive characterization and the corrosion-only and tribocorrosion tests conducted in this study, the following conclusions can be drawn:
Results indicated that 316AP had an Eb higher than those of 316W and 317W. Following annealing, the Eb of SLM 316L decreased to values close to those of 317W.
The additively manufactured 316L stainless steel had an abundance submicrometer spherical Si- and O-rich amorphous inclusions in their microstructure, as revealed by TEM-EDS. It was found that solution annealing resulted in the formation of larger crystalline oxide inclusions.
NanoSIMS results indicated that the passivity of the wrought 317W and 316W was found to be similar, despite the differences in their Mo level. However, under tribocorrosive conditions, 317W exhibited the most stable passive behavior, even at extreme anodic potentials of up to 1.0 VSSE. These findings were confirmed by both polarization tests and the interrupted potentiostatic results.
Under tribocorrosive conditions, the main difference between SLM fabricated and wrought samples appeared at applied anodic potentials greater than +0.2 VSSE.
In contrast to 316W and 317W, the SLM-fabricated samples showed dissimilar chemical and mechanical breakdown behaviors. The PPM tests indicated that SLM-fabricated samples repassivated at a similar rate to that of 317W. However, following tribocorrosion tests, the SLM-fabricated samples had a poorer repassivation behavior than that of 317W, attributed to the higher Mo content of 317W.
UNS numbers are listed in Metals & Alloys in the Unified Numbering System, published by the Society of Automotive Engineers (SAE International) and cosponsored by ASTM International.