Corrosion Properties of Powder Bed Fusion Additively Manufactured 17-4 PH Stainless Steel

Metal additive manufacturing (AM) has recently become a desirable process for complex parts across a broad range of applications.¹  Laser powder bed fusion (LPBF) is an AM process for metals whereby a laser is used to selectively melt a pattern in successive layers of powder material as a means of building a three-dimensional structure. The locally high cooling rates produced during the process with highly non-equilibrium solidification conditions result in microstructures that can vary significantly from traditional wrought materials.^2-3  LPBF processes can create material that often contains substantial solute segregation with formation of terminal solidification phases. It can also contain unprocessed particles from the starting powder, resulting in pores.^2-3  As a consequence of these varied microstructures, AM alloys may exhibit properties vastly different from their conventional wrought or cast counterparts.^4-5  For example, Yadollahi reported that fatigue cycles for a range of stress amplitudes for 17-4 PH AM in the H1050 condition were nearly an order of magnitude less than wrought material. This was ascribed to the large number of defects in the AM material, including pores, unmelted regions, and unmelted powder particles.⁶  Mower, et al., found that AM 17-4 PH displayed a considerably diminished yield strength compared to wrought material (610 MPa to 737 MPa vs. 898 MPa) and attributed this to incomplete fusion across build planes in the AM process.⁷  While some adjustments have been made in manufacturing to enhance mechanical traits, such as incorporation of carbide particles into the powder mixture for strengthening, very little attention has been directed at understanding or controlling the corrosion properties of these materials.^8-9 

In this study the authors explore the relative impact of microstructural characteristics unique to AM processing on the corrosion properties of LPBF 17-4 PH, a precipitation hardened, martensitic stainless steel (UNS S17400)⁽¹⁾. Specifically, and for the first time, the corrosion behavior of a 17-4 PH LPBF material under full immersion saline conditions is investigated and comparison to conventional wrought 17-4 PH stainless steel is provided. Studies on 300 series AM stainless steels suggest that porosity of these materials can considerably increase corrosion susceptibility, but direct evidence of the influence of pores on corrosion properties is lacking.^3,10  Here both bulk and local electrochemical measurements are utilized to relate microstructural features to corrosion behavior, including the role of pores on passivity. The results presented provide important feedback for enhancement of AM processing techniques to optimize corrosion properties.

Studies to date of LPBF manufactured stainless steels have uncovered substantially higher corrosion susceptibility compared to wrought counterparts. It has been suggested that elemental segregation, residual oxides from initial powders, and porosity, similar to characteristics seen in more conventional powder metallurgy (PM) materials, may be at fault.^3,11-12  AM materials are expected to exhibit reduced corrosion properties similar to conventional PM materials, due to the comparable microstructures formed from processing.¹² 

In PM stainless steels, elemental segregation from non-equilibrium processing conditions reduces the corrosion resistance of the materials compared to conventional wrought or cast equivalents. Chromium depletion of stainless steel is of particular concern, as it results in regions with reduced or unstable oxides that can lead to locally sensitive areas prone to corrosion attack.¹³  Samal and Terrell demonstrated that hydrogen sintered 316 stainless steel (UNS S31600) processed with a range of cooling rates and carbon content (0.01% to 0.11% C) resulted in a range of sensitized material. The authors applied a critical cooling temperature curve for wrought 316L (UNS S31603) to avoid sensitization, showing a consistent correlation between sensitization and corrosion rates for the 316L PM samples.¹⁴ 

Initial studies of AM processed stainless steel exhibit trends similar to PM materials regarding corrosion resistance and elemental segregation of passivating elements. Trelewicz, et al., found that microsegregation in AM materials degrades their corrosion properties. Segregation of Mo in 316L due to LPBF processing increased the passive current density of the AM material by almost an order of magnitude in 0.1 M HCl solution.¹¹ 

Elevated oxygen content in PM alloys originating from starting powders is known to cause a considerable decrease in corrosion resistance, a highly probable scenario for AM materials as well. Oxide formation originates from relatively large amounts of pre-existing oxides on the high surface area starting powder. These oxides are not always reduced in the furnaces and can lead to a reduction in the pitting potential of PM steels.^15-17  Klar, et al., reported a strong dependence of 316L corrosion initiation time (with visible surface staining or rust on the surface) in 5 wt% NaCl electrolyte on oxygen content, whereby initiation time was near 10⁴ h for 250 ppm oxygen and decreased to less than 0.5 h for 1,750 ppm.^16-17 

The inherently porous nature of PM stainless steels has been demonstrated to play a governing role in their reduced passivity.^12,18-19  Pore geometry and interconnectivity are two critical factors that can govern occluded cell-type corrosion susceptibility. Pores breaching the material surface with favorable attributes in this regard can serve as crevice zones wherein IR drop and maintenance of acidic, hydrolytic conditions can enable depassivation within the pores. One study on 304L (UNS S30403) and 316L exposed to ferric chloride solutions resulted in primary corrosion attack occurring within pores.¹⁹  Similar results on 304L and 316L were seen with sulfuric and phosphoric acids. The authors attributed this attack to local acidification within the pore, resulting in depassivation of the material and formation of an active region within the pore, supported by cathodic activity outside the pore.¹⁸ 

Investigators have demonstrated general trends between porosity and corrosion resistance, but this is complicated by the aforementioned factors and the environment. Jones reported that corrosion resistance of three austenitic stainless steels with varying wt% Ni improved with decreased porosity for a sintered density range of 6.25 g/cm³ to 7.25 g/cm³ in acid medium, 40% HNO₃.²⁰  However, in saline solutions, the relationship between material density and corrosion susceptibility was complicated further by pore geometry on 316L. At high porosity and large pore size ≥20 μm, little to no crevice corrosion was promoted. At intermediate porosity, high corrosion occurred as the pores are narrow (≤20 μm), promoting hydrolysis and leading to crevice corrosion, and at high density, few pores existed to promote crevice corrosion.^21-22  A similar relationship between corrosion rate and pore size for 316L was shown for atmospheric salt fog exposures, where corrosion rate was inversely related to pore diameter.²³ 

Examinations have shown that localized corrosion attack on AM materials can also be initiated by defects such as elongated grains, heterogeneities,¹¹  and porosity.¹⁰  Porosity and surface roughness were shown to play large roles in the corrosion properties of AlMgSi direct metal laser sintered alloys, where even in the polished condition, remaining pores exhibited enhanced corrosion attack in SEM micrographs.¹⁰  While these studies provide initial insight into the corrosion behavior of AM materials, the literature lacks results relating direct, local scale exploration of either the chemical or the morphological microstructural characteristics of AM materials on the corrosion properties of the material.

In this study the properties of wrought and LPBF 17-4 PH in the H900 temper condition are explored via full immersion and local microelectrochemical measurements in 0.6 M NaCl solution to establish a baseline. Mechanical properties of samples from this same material build have been studied and published elsewhere.²⁴  Correlation is drawn between the measured corrosion behavior and microstructural features characterized by optical and electron microscopy techniques before and after electrochemical interrogation. The results provide direct evidence of the primary governance of pores on the decreased corrosion resistance of the LPBF material. The mechanisms behind the effect of porosity on corrosion behavior, future investigations necessary for enhanced understanding, and possible solutions/suggestions for AM processing to enhance the corrosion resistance of these materials are discussed.

Additively manufactured samples of 17-4 PH stainless steel were compared to conventional wrought samples. The AM materials were acquired commercially from an outside vendor that utilized a ConceptLaser Mlab^† printer. The AM samples were printed by a laser powder bed fusion process, in which feedstock powder was sintered by a laser in successive layers, with preprogrammed factory settings for patterning. All LPBF 17-4PH specimens were heat treated to industrial material standards (AMS 5604) by solution heat treating at 1,050°C for 60 min in argon atmosphere and subsequently cooled to room temperature. Both AM and wrought specimens were then age hardened to the H900 condition by heating to 482°C for 60 min in air.

All sample surfaces tested or examined were ground to 1200 grit silicon carbide paper, polished to 1 μm diamond paste, cleaned with 18.2 MΩ deionized (DI) water, then isopropyl alcohol, and dried with nitrogen. The top surface was selected for study, as it was the surface available in a testable sample size; see Figure 1. Samples of the wrought material were prepared in the same manner. Polished LPBF samples for bulk electrochemical experiments were tested both pre- and post-sonication in DI water, and no significant difference was observed. However, materials for microscale experiments were sonicated to avoid the influence of possible trapped polishing particles.

Material characterization of the LPBF and wrought 17-4 PH was performed to determine both composition and microstructure. Compositional analysis of the as-received materials was determined via inductively coupled plasma atomic emission spectroscopy (ICP-AES) and ICP-mass spectrometry (MS) for Cr, Ni, Mn, Cu, Si, Mo, P, Cb, and Ta and via LECO^† furnace analysis for C, S, O, and N. Fe was determined by the difference in these measurements. Microstructural analysis comprised standard optical metallurgical examination along with scanning electron microscopy (SEM), energy dispersive spectroscopy (EDS), and wavelength dispersive x-ray spectroscopy (WDS). Light optical microscopy was performed on polished sections etched using Viella’s reagent, which is a solution composed of 100 mL ethanol, 1 g picric acid, and 5 mL concentrated hydrochloric acid. The samples were etched at room temperature (23°C) for 60 s. As-polished samples were utilized for other analyses. SEM/EDS were performed at 20 keV at high vacuum mode and at a working distance of 9 mm to 12 mm, using a Zeiss^† field emission source SEM. EDS signals were normalized to the maximum peak (Fe K-α). Electron probe microanalysis (EPMA) via WDS was performed using a JEOL JXA-8530F HyperProbe Electron Probe Microanalyzer^†. High-resolution beam scan maps were acquired using an accelerating voltage of 15 keV with a beam current of 20 nA using a point-to-point spacing of 0.1 μm. A ZAF correction procedure was utilized to produce quantitative elemental maps of the LPBF 17-4PH microstructure.

Pore number density and size distribution in the build direction plane and perpendicular to the build direction plane, shown in Figure 1, were estimated via image analysis of secondary electron micrographs and optical images. ImageJ^† software was applied for this analysis, using a threshold procedure.²⁵  A minimum of three images were analyzed for each direction, providing pore distribution and size information on the build direction plane and perpendicular plane.

Electrochemical experiments were performed using either a standard three-electrode flat cell to characterize bulk behavior or a microelectrochemical three-electrode cell to target the behavior of specific surface features. The electrolyte used for all experiments was ambiently aerated, pH 6, 0.6 M NaCl solution prepared with 18.2 MΩ deionized water. All electrochemical experiments were evaluated at room temperature (23°C).

Both open-circuit potential (OCP) and potentiodynamic polarizations were performed in the flat cell. A saturated calomel electrode was used as the reference electrode and PtNb mesh as the counter electrode. Prior to potentiodynamic measurement, samples were held at OCP for 1 h or 24 h to allow materials to come to steady state. Anodic and cathodic potentiodynamic scans were conducted at a scan rate of 0.1667 mV/s after the open-circuit holds. A minimum of five replicate scans for each material and condition were conducted.

A microelectrochemical cell method following procedures outlined elsewhere was used to locally target surfaces with different pore size populations.^26-32  The microelectrochemical cell used consisted of a silicone-tipped glass capillary with an inner diameter of 380 μm connected to a cell body containing a Pt wire as a counter electrode and chloridized silver wire as a reference. The chloridized silver wires were made following procedures developed by Hassel, et al., and allowed to stabilize in 0.6 M NaCl for 12 h.³³  Measurements were done by touching down the cell filled with 0.6 M NaCl solution to targeted areas on the LPBF and wrought samples to create a seal. OCP measurements were then taken for 1 h and 24 h along with subsequent potentiodynamic measurements at a scan rate of 1 mV/s. A faster scan rate was selected for the microelectrochemical cell than that used in full immersion experiments in order to avoid H₂ bubbles during initial cathodic polarization that could block the capillary and to minimize changes in solution chemistry in the capillary. No crevice corrosion was observed after electrochemical testing. The microelectrochemical cell aids in establishing the influence of pore size populations, large vs. small, on the corrosion behavior of the LPBF material.

Polished AM samples were immersed in ambiently aerated 0.6 M NaCl solution up to 7 d. After exposure, samples were rinsed with deionized water, sonicated for a minimum of 5 min to remove excess salt solution, and dried with nitrogen prior to imaging in the SEM.

Compositions of the LPBF and wrought 17-4 PH studied here are presented in Table 1. Figure 2 is a micrograph showing the typical appearance of the etched microstructure of wrought and AM 17-4PH samples. This material can be characterized as a martensitic matrix with a relatively fine (∼10 μm to 20 μm) prior austenite grain size, with interspersed islands of delta-ferrite. No evidence of solidification substructure can be observed in these micrographs.

Pore distribution and microstructural morphology, resulting from nonequilibrium solidification conditions of AM processing of the LPBF material, are shown in SEM micrographs in Figures 3 and 4 and in the EPMA images in Figure 5. Larger pores (diameter [d] > 50 μm) exhibit very irregular shapes and can contain remaining unsolidified particles, whereas smaller pores (d < 10 μm) are typically more hemispherical in shape (noted in Figure 4[a]). These smaller spherical pores have been observed in other AM materials and can be attributed to gas porosity rather than unmelted regions within the LPBF sample.^1,34  Figure 4(a) shows a higher-magnification image of a typical pore in which unsolidified powder particles from the printing process can be seen. An EDS scan was taken of the entire area, displaying even Cr distribution across the entire surface (Figure 4[b]; within the pore, line of sight to the EDS detector limits some detection of Cr). Three spectra are shown for comparison, from the undissolved particle, the inner surface of the pore, and the general polished surface of the LPBF material, showing no considerable gross difference in elemental composition among these three regions. High-resolution EPMA maps of the LPBF specimens after heat treatment (Figure 5) show a lack of discernable microsegregation in the martensitic matrix (again, within the pore, line of sight to the WDS detector limits some detection of Cr). Local enrichment of Nb and Si/O corresponding to Nb-carbonitrides and Si-rich oxides is observed.

Measurements of the maximum pore diameter on the top surface and a cross section of the 17-4 PH AM material (Figure 1) reveal similar populations in both directions (Figure 6). The average diameter was found to be 14 μm for the top surface, ranging from 3 μm to 273 μm, and 13 μm for the cross-sectional surface, ranging from 3 μm to 295 μm (the smallest detectable features were in the range of 2 μm to 4 μm at the magnification taken in the SEM). The percent area coverage was 3±0.7% for the top surface and 2.4±0.2% for the cross-sectional surface. For all electrochemical testing, the top surface, parallel to the print plane, was examined as it was the readily available surface in the dimensions necessary for testing. However, as the distribution and cross-sectional geometry of the pores seen on the cross-sectional surface did not differ greatly from the top surface, similar electrochemical results would be expected.

Representative OCP vs. time for 24 h exposures are plotted in Figure 7. The OCP of the wrought material rises gradually with time. However, for the LPBF material, the OCP changes erratically over the exposure time and appears to stabilize after about 19 h of exposure. In addition, the OCP of the LPBF sample is lower than that of the wrought material.

The average OCP from the 1 h and 24 h scans was calculated and then compared to typical values from the literature for wrought 17-4 PH (Table 2). Values for the wrought sample tested at 1 h are comparable (within 30 mV) to those found in the literature for 17-4 PH stainless steels tested for 15 min to 40 h.

Representative anodic polarizations after 1 h at OCP displayed different behavior between wrought and LPBF 17-4 PH. The wrought material exhibited spontaneous passivity, whereas the LPBF material exhibited higher anodic currents with no region of apparent passivity. Current densities were approximately two orders of magnitude greater over the same voltage range (Figure 8). However, for the anodic scans taken post 24 h OCP, both the LPBF sample and the wrought sample exhibited passive regions. Again, the current densities were higher by about two orders of magnitude. Cathodic scans for both wrought and LPBF samples post 1 h OCP displayed similar behaviors, with a lowered OCP for the LPBF material (Figure 9), indicating a similar behavior in cathodic kinetics for the surfaces of both samples.

A comparison of the open-circuit corrosion current densities estimated for each material (i_corr) for the LPBF vs. wrought 17-4 PH samples is given in Table 2. These were calculated using the Tafel slope extrapolation, however, they are meant as indicators rather than true corrosion current densities, as 17-4 PH exhibits passive behavior with pitting corrosion rather than general corrosion. It is interesting to note that the LPBF samples demonstrate a higher corrosion current density, by nearly an order of magnitude, for the anodic scans taken post 1 h and 24 h holds at the OCPs.

The results of the microelectrochemical cell analysis for the 1 h OCP are plotted in Figure 10(a) and display no considerable differences between the large-pore (d ≥ 50 μm) and small-pore (d ≤ 10 μm) surfaces. Small pores are included in both scan measurements, as they are distributed fairly evenly across the LPBF surface and were unavoidable in the current measurement technique. However, as can be seen in Figure 10(b), the anodic scan taken above the large-pore surface showed increased current densities with no apparent passive region and a lower OCP than the small-pore surface. The scan above the small pore demonstrated a passive region, more similar to that of the wrought material. A microelectrochemical scan of a wrought sample is also plotted in Figure 10(b). A comparison of i_corr values for the LPBF 17-4 PH large pores, small-pore surfaces, and wrought material is given in Table 3. These are again meant only as an indication of local passive current densities. In this case, the current densities are much closer, and no significant difference is seen between the i_corr of the large-pore and small-pore surfaces. As a comparison, OCP holds were taken for 24 h followed by anodic scans and are plotted in Figure 11. As can be seen, the OCP for the larger pores is lower and less stable than that for the smaller pores. However, both anodic scans display a large passive region similar to that of the wrought material (Figure 11[b]).

Post-exposure, for samples immersed in 0.6 M NaCl solution for 1 week, SEM imaging was applied and images for an LPBF sample are shown in Figure 12. Corrosion product can be seen forming near and around pores or possible pits. While these are smaller-scale pores, d ≈ 10 μm to 20 μm, they still exhibit irregular shapes with crevice-like regions due to lack of fusion during processing.

The LPBF 17-4 PH material examined in this study displayed inferior corrosion resistance, with increased passive current density and decreased passive region compared to wrought 17-4 PH materials in the same environment. Previous studies of PM materials bearing similar microstructural and morphological features have attributed this to elemental segregation, retained oxides, and/or porosity, as discussed earlier. In this study, the presence of pores was found to be the primary contributing factor to corrosion susceptibility. The following discussion rationalizes this finding.

Material composition analysis given in Table 1 shows that the LPBF samples are within the accepted typical requirements for 17-4 PH. Conventional wrought 17-4 PH is a precipitation-hardened stainless steel with 15 wt% to 17.5 wt% Cr. At this elemental composition, this alloy is close to the lower limiting composition of what is considered a stainless steel, 11 wt% Cr.³⁵  The amount of Cr is limited in this alloy to allow for a fully transformed martensitic structure at room temperature. Simple, empirically derived relationships originally proposed by Eichelmann and Hull were used to relate alloy composition and the start temperature of the martensite transformation.³⁶  For the alloys examined here, all LPBF samples possessed an estimated M_s temperature between 265°C and 285°C. For comparison, the estimated M_s temperature for wrought 17-4PH sheet was 248°C. While these are empirical estimates, such calculated transformation temperatures are indicative of the precondition of austenite instability at room temperature for 17-4PH. However, as the Cr composition of 17-4 PH is close to passivity limits, any changes to the microstructure, in either elemental segregation, oxide formation, or morphology, may affect the corrosion susceptibility of the material, as they could easily disrupt the passivating oxide layer.

The melting and rapid solidification of 17-4PH during laser powder bed fusion processes results in microstructural features distinct from thermomechanically processed material such as sheet. Characteristic features from the solidification process include relatively large solidification grains that grow preferentially along an axis aligned toward the moving heat source. Moreover, solidification grains will also contain substructure that is formed by operative solute redistribution during nonequilibrium solidification.³⁷  For precipitation-strengthened martensitic stainless steels, including LPBF 17-4PH, heat treatment of the as-fabricated material is required to produce the desired mechanical and corrosion behavior. Figure 2 shows light optical micrographs of etched AM 17-4PH microstructures for the corrosion tests examined in this work. The complete lack of solidification features in the heat-treated LPBF 17-4PH samples indicates that the 1,050°C solutionization heat treatment resulted in recrystallization. The heat treatment also led to homogenization of solute microsegregation, and the subsequent microstructure was lacking in obvious microscale features expected to lead to gross reduction in passivity. Both EDS and WDS elemental maps exhibited elemental uniformity across the samples (Figures 4 and 5). It has already been seen that Cr depletion around Cr carbides that precipitate in these conventional materials during hardening can result in a more active pitting potential by 20 mV to 60 mV in 3.5 wt% NaCl solution.³⁵  Furthermore, as the C and N compositions were 0.017 wt% and 0.036 wt%, respectively, no chromium carbides or chromium nitrides associated with Cr depletion and sensitization were observed through WDS.

While PH martensitic stainless steels typically have improved corrosion properties compared to martensitic stainless steels,³⁸  the precipitation-hardening process can slightly reduce the corrosion resistance of these steels due to the formation of precipitates.³⁵  Elevated amounts of niobium carbonitrides and silicon oxides relative to conventional wrought material were present in the LPBF samples. These were small and evenly distributed throughout the surface and not localized at the pores (Figure 5). Niobium carbonitrides have been found to be beneficial to the corrosion resistance of stainless steels, as they suppress the formation of chromium carbides.³⁹  With respect to silicon oxides, it has been previously established that increased amounts of Si can actually enhance corrosion resistance in stainless steel alloys.^40-41  However, Castle, et al., have also found that pitting can occur at corrosively inert oxides with some dissolution of the oxide.⁴²  While silicon oxide particles in the LPBF material could contribute to enhanced corrosion, they are not believed to be a dominant microstructural factor, as evidenced by the microeletrochemical cell experiments.

The primary difference found between the LPBF samples and conventional wrought material is the high porosity seen in the materials in Figures 3 and 6. Two distinct morphologies were observed: large, irregularly shaped pores (d > 50 μm) due to lack of fusion defects (Figure 4[a]) and smaller, spherical pores (d < 10 μm) formed by gas entrapment and coalescence during processing (Figure 5[a]). The influence of these structures on the corrosion behavior of the LPBF material is evident in the electrochemical results.

Electrochemical measurements of the OCP provide insight into the origin of decreased passivity of the LPBF samples, exhibiting lower and more unstable OCPs than the wrought material (Tables 2 and 3, and Figure 7). This behavior is consistent with previous work on porous powder metallurgy samples relative to their wrought counterparts.^43-45  Instabilities in the OCP of the LPBF samples could be due to solution ingress into the pores or active corrosion/pitting (Figures 7 and 11).^45-46  The overall decreased potential on the LPBF sample could be a sign of thinning or localized attack of the passive film, whereas the rise in OCP of the wrought material is indicative of an increase in passive film thickness on the material over time.⁴⁶  Another possibility for this change in potential is related to the polarizability of stainless steel, in that the surface potential could be easily affected by localized corrosion events occurring across the surface.

Anodic polarization scans revealed reduced corrosion resistance for the LPBF samples as compared to the wrought material, but also a dependence on pre-exposure time. It is interesting to note that the LPBF samples pre-exposed to a 24 h OCP displayed a passive region in the anodic scan (Figure 8[b]), while those pre-exposed for only 1 h at OCP did not (Figure 8[a]). One possible explanation for this phenomenon could be that the conditions within the pores change significantly such that they do not meet critical crevice conditions through either growth of the pore that changes the critical geometry or a change in solution chemistry.⁴⁷  Thus these sites may deactivate as regions for enhanced anodic dissolution and the material exhibits a passive region more similar to that of the wrought material. This is evident by the anodic scans of large and small pores after a 24 h OCP hold, shown in Figure 11. The OCP holds still display a lower, more unstable OCP over the 24 h period for the larger pore sample (Figure 11[a]). However, the anodic scans post 24 h OCP hold result in a larger passive region for the large pore sample, indicating passivation of the pore surface with time (Figure 11[b]).

The additional LPBF surface area imparted by porosity likely does not account for the 5× to 10× higher open-circuit corrosion currents estimated for the LPBF samples relative to wrought (Table 2). Assuming the LPBF pores are hemispherical and using the measured pore size and distribution in Figure 6, it can be estimated that their presence would only increase the surface area by a factor of 1.06. This is insufficient to account for the disparities in the current density measured. This is further verified by cathodic polarization scans (Figure 9) displaying the same current density for both samples, indicating that measurements were made from the same surface area. It must be recognized, however, that these larger-diameter pores are not hemispherical, nor do they have a smooth round surface (Figures 3 and 4). Results from microelectrochemical cell experiments provide further insight into these enhanced corrosion rates.

Microelectrochemical cell experiments established that pore features directly affect the passivity of the AM stainless steel. In Figure 10, anodic polarizations over areas with large pores (d > 50 μm) exhibit active corrosion behavior with increased current density over the same voltage range. This is in contrast to passive behavior comparable to that of the wrought material witnessed in areas with few small pores (d < 10 μm). The average breakdown potentials of the large pores are >400 mV below that of the wrought material, 311±27 mV_SHE for the large pores compared to 743±5 mV_SHE for the wrought material. Breakdown potentials at small pores are more comparable to the wrought material at 577±70 mV_SHE. This behavior changed after a 24 h OCP hold. Anodic scans on large pores exhibited a passive region with still a lower OCP, suggesting that passivation of the pore surface occurred during the OCP hold. The enhanced corrosion behavior of the LPBF 17-4PH material is primarily due to the presence of large pores at the sample surface. This effect may be related to pore geometry, where crevice corrosion can initiate in large pores, as they tend to have more narrow crevice-like geometries due to unmelted powder particles vs. smaller hemispherical pores formed by gas coalescence. Further indication of enhanced corrosion at pores was seen post-exposure in full immersion, where corrosion appeared at or near pores (Figure 12).

As pores can act as occluded areas that deplete in oxygen and become acidified during initial stages of corrosion, they can be even more conducive to corrosion attack.⁴⁸  Mathiesen, et al., observed a geometrical dependence of corrosion on PM materials in that pores of a smaller diameter (d) had a higher corrosion severity.⁴⁹  An empirical relationship was previously developed for atmospheric salt fog exposures of PM 316L, in which corrosion severity was found to be inversely proportional to the pore diameter.²³  This suggests that as pore diameter decreases, corrosion severity increases. However, this does not account for any limit to enhancement of the severity as d tends to zero for the pores, which would be expected, as in that case the material should resemble a conventional wrought alloy. This relation also lacks a parameter that would provide information about the aspect ratio of the pore. Pore geometry relationships can be compared to traditional crevice corrosion geometry relationships; Oldfield and Sutton found a strong dependence of crevice pH on crevice depth, thus deeper crevices could establish more severe environments for corrosion attack.⁵⁰  The interplay of pore diameter and crevice solution can affect the depth of the active/passive region, where the depth of attack increases to a limit for a wider crevice gap and a specific crevice depth.⁵¹  Others have established a scaling law between the geometry of a crevice and corrosion susceptibility, where the relationship between maximum crevice attack, x_crit, and the crevice gap, G, is either or x_crit/G, depending on the passivity of the crevice walls and whether it was IR-controlled crevice corrosion.^52-54  Therefore, both pore diameter and depth can affect corrosion attack.

In the current work, larger pores were seen to exhibit a more reduced passive region and enhanced corrosion currents compared to the smaller pores in the microelectrochemical experiments. This may have to do more with pore geometry than size. Larger pores, formed from lack of fusion, appeared more varied in shape, with regions that could be considered tight crevices, as seen in Figure 4(a). Smaller pores were generally formed due to gas coalescence and were hemispherical in shape, lacking the tight crevices that drive the occluded cell corrosion (Figure 5[a]). More detailed investigation of the relationship between pore geometry and its effect on corrosion rate would aid in understanding the influence of pore size, morphology, and density on corrosion susceptibility in LPBF materials. Possible application of crevice modeling to pore geometries in AM materials could aid in understanding the severity of pores in terms of corrosion resistance.

As this study has found porosity to be the primary contributing factor for enhanced corrosion of AM materials, possible solutions for reducing corrosion susceptibility should include processing to reduce surface porosity. Solutions for decreasing the corrosion susceptibility of PM materials with respect to surface roughness have been previously explored. Grinding, turning, or shot-blasting surfaces can seal surface pores and improve their corrosion resistance.⁵⁵  Thermal and/or chemical passivation processes can alter the oxide layers and improve corrosion resistance.^55-56  These or similar techniques could be applied to AM materials to alleviate the negative effects of porosity on corrosion.

Through electrochemical measurements and immersion exposures, it was found that LPBF 17-4 PH exhibited reduced corrosion resistance compared to conventional wrought 17-4 PH material in 0.6 M NaCl. Porosity was identified as the primary cause of the decreased corrosion resistance.

• Microstructural evaluation of the LPBF 17-4 PH material revealed porosity to be an obvious differentiating feature of consequence to passivity compared to wrought 17-4 PH samples. Cr segregation was not believed to greatly affect the corrosion properties, as heat treatments were performed for homogenization of solutes. Oxides in the LPBF material may play a role but were not the dominant factor in 17-4 PH.

• LPBF vs. conventional wrought material exhibited decreased corrosion resistance through electrochemical testing. Enhanced corrosion rates on the LPBF samples over the wrought samples were indicated by the calculated i_corr values for the full-immersion electrochemical experiments.

• Microelectrochemical cell experiments established that the presence of pores on the sample surface directly affects the corrosion type, exhibiting active corrosion above the large pores (d ≥ 50 μm) rather than the passive behavior displayed above regions with smaller pores (d≤10 μm) and on wrought material.

• Further evidence of enhanced corrosion at pores was confirmed by post-exposure analysis of full immersion exposures, where corrosion appeared to initiate at or near pores.

• Further understanding of the influence of porosity, such as pore size and aspect ratio, on the corrosion properties of LPBF stainless steels could help to optimize processing parameters or post-processing procedures to enhance corrosion resistance. For the material under study here, a decrease in the larger range of porosity resulting from lack of powder fusion, through either processing or post-processing treatments, would be expected to dramatically enhance corrosion resistance.

The authors gratefully acknowledge Bonnie McKenzie for SEM and EDS measurements and analysis, Alice Kilgo for metallographic sample preparation, and Dick Grant for EPMA and WDS measurements. Sandia National Laboratories is a multiprogram laboratory managed and operated by Sandia Corporation, a wholly owned subsidiary of Lockheed Martin Corporation, for the U.S. Department of Energy’s National Nuclear Security Administration under contract DE-AC04-94AL85000.