As a high-strength corrosion-resistant alloy, stress corrosion cracking (SCC) behavior is a key consideration for the conventional, wrought form of 17-4PH stainless steel. With the increasing popularity of the additively manufactured (AM) form of 17-4PH, understanding the SCC behavior of AM 17-4PH will be similarly critical for its presumed, future applications. The current study quantifies and compares the SCC behavior of both the wrought form, as a baseline, and AM form of 17-4PH at peak-aged (∼1,200 MPa) and overaged (∼1,050 MPa) strength levels. The laser powder bed fusion technique followed by post-process hot isostatic press (HIP), solution annealing, and aging heat treatments is used to produce AM 17-4PH with similar microstructures and strength levels to wrought 17-4PH and facilitate the comparison. SCC behavior is quantified using fracture mechanics-based rising (dK/dt = 2 MPa√m/h) and constant (dK/dt = 0 MPa√m/h) stress intensity tests in neutral 0.6 M NaCl at various applied potentials. Limited SCC susceptibility was observed at open-circuit and anodic potentials for both forms of 17-4PH. At cathodic applied potentials, AM consistently underperforms wrought with up to 5-fold faster crack growth rates and 200 mV to 400 mV wider SCC susceptibility ranges. These results are interrogated through microstructural and fractographic analysis and interpreted through a decohesion-based hydrogen-assisted crack model. Initial analyses show that (1) increased oxygen content, (2) porosity induced by argon processing, and (3) slow cooling (310°C/h) during conventional HIP processing might contribute to degraded SCC performance in AM 17-4PH.

The precipitation-hardened, martensitic stainless steel 17-4PH (UNS S17400(1)) exhibits both high strength and good corrosion resistance.1-2  Due to these favorable properties, it is widely used in aggressive operating environments across the marine, petrochemical, nuclear, and aerospace sectors. While 17-4PH has a long history in its conventional wrought form, a new form of 17-4PH is becoming popular through the use of additive manufacturing (AM) technologies. In particular, the laser powder bed fusion (LPBF) technique3  with powder meeting the 17-4PH composition is being used to generate “additive” 17-4PH. Unlike wrought 17-4PH, which exhibits a martensitic microstructure, the rapid solidification process associated with LPBF typically results in a reduced-strength ferritic microstructure in the as-built condition.4-5  However, the underlying composition of AM 17-4PH gives it the potential to reach the strength of its wrought counterpart through appropriate heat treatment. The strength of heat-treated AM 17-4PH can meet or exceed the strength of wrought 17-4PH.6-8 

Traditionally manufactured 17-4PH can be susceptible to environmental degradation; specifically, in-service failures have been reported for 17-4PH marine vessel propeller shafts,9  fasteners on spacecraft,10-11  valves on petrochemical pipelines,12-13  various oilfield equipment,1,14-15  nuclear reactor components,16  among others.2,17  In many of the above instances,9-13  component failure was attributed to hydrogen-driven stress corrosion cracking (SCC) based on the observation of brittle fracture morphologies and the presence of a known hydrogen-rich environment. Laboratory testing has established that the susceptibility of wrought 17-4PH to SCC is strongly dependent on environmental conditions.16,18  For example, high resistance to SCC was observed at open-circuit potentials (OCP) in dilute chloride solutions,16,19-23  while increased SCC susceptibility was noted when 17-4PH was exposed to cathodic polarizations in similar environments.16,20-21,24-25  Additionally, a strong heat treatment dependence on the SCC behavior was observed in dilute chloride environments, with peak-aged alloys often exhibiting increased susceptibility to SCC relative to solution-treated or overaged conditions.11,16,20-21,2527  Interestingly, significant heat-to-heat variability was also noted during salt spray-based SCC experiments, even in the overaged condition.28 

The clear dependence of SCC behavior in wrought 17-4PH on heat treatment and the reported heat-to-heat variability strongly suggest an important role of the underlying microstructure in establishing the SCC susceptibility of 17-4PH. As such, it is likely that AM 17-4PH will exhibit different SCC behavior relative to a similar strength wrought 17-4PH alloy as important differences in microstructure are expected. For example, though properly heat-treated AM 17-4PH will have the traditional precipitation hardened martensitic structure, it is likely that the ferritic microstructure of the as-built condition will induce a different austenite morphology during the solution-annealing step as compared to wrought 17-4PH, thus impacting the morphological details of the martensitic structure. Additionally, AM 17-4PH may contain specific defects that will not be present in wrought materials, such as porosity, lack of fusion defects, and delamination between build layers.29  AM alloys have also been reported to contain unique secondary phases such as oxide inclusions30-33  and have larger prior austenite grain (PAG) sizes.34  To date, studies of AM 17-4PH has focused on microstructure, strength, fatigue,29,35-38  and general corrosion behavior.39-41  However, despite the importance of the SCC failure mode for 17-4PH components and the likelihood of differences in SCC behavior for AM and wrought 17-4PH, studies of SCC behavior are limited.34 

The purpose of this study is to quantify and compare the SCC performance of AM and wrought 17-4PH at similar strength levels and identify the microstructural features responsible for any differences. First, the SCC behavior as a function of applied potential under full immersion in 0.6 M NaCl solution is measured for both AM and wrought 17-4PH under various heat-treatment conditions using fracture mechanics-based testing. High-magnification microscopy and electron backscatter diffraction (EBSD) are then used to identify key differences in microstructural parameters. The mechanistic influence of observed microstructural variations on the SCC behavior is interpreted in the context of an established decohesion-based model for hydrogen-driven SCC and the implications of the current findings are then discussed.

Materials

Wrought and AM 17-4PH were obtained for comparison in this study. The AM material was built using the LPBF process with an EOS M290 machine. In accordance with best practices, the built-in machine parameters recommended by EOS for 17-4PH were used.42  Notable parameters include a 47° hatch rotation angle, 0.11 mm infill hatch distance, 220.1 W infill laser power, 755.5 mm/s infill laser speed, 0.04 mm layer thickness, 80°C build plate temperature, an argon build atmosphere, and the use of argon-atomized powder (particle size range of 15 µm to 45 μm). Two LPBF builds were performed according to the layout in Figure 1 to produce 21 mm × 21 mm × 75 mm stock bars in two orientations. Note that the coordinate system shown conforms to ISO/ASTM 52921,43  and the stocks bars were rotated 11° off axis to aid with powder flow during recoating. Virgin powder was used for the first build and then recycled for use in the second build.
FIGURE 1.

LPBF build layout used to generate additive stock material. Various specimen geometries (SENT, tension, excess) were machined postbuild from the two stock orientations (ZXY, XYZ) as shown.

FIGURE 1.

LPBF build layout used to generate additive stock material. Various specimen geometries (SENT, tension, excess) were machined postbuild from the two stock orientations (ZXY, XYZ) as shown.

Close modal

Postbuild, the printed bars of stock material were cut from the build plate using electrical discharge machining (EDM). A hot isostatic press (HIP) and then standard44  solution annealing (SA) heat treatment was performed on the stock. The HIP treatment was performed in argon at a pressure of 100 MPa and temperature of 1,125°C for 4 h, followed by furnace-cooling at 310°C/h to 300°C and then air cooling to ambient temperature. The HIP time, temperature, and pressure were chosen based on recommendations in ASTM F330145  for 4340 and UNS S31603 because specific parameters for 17-4PH were not available. The HIP cooling rate was the standard rate used by the HIP process supplier. The SA treatment was conducted in lab air at 1,040°C for 0.5 h, followed by air-cooling to ambient temperature. After the high-temperature heat treatments, single edge notch tension (SENT), tension, and excess specimens were machined from the stock (Figure 1). Each machined specimen underwent an aging heat treatment in lab air for either peak strength (peak age; 460°C for 1 h, then air-cooled) or increased toughness (overage; 552°C for 4 h, then air-cooled). Note that the AM peak aging temperature differs from those conventionally used for wrought 17-4PH based on EOS recommendations42  for reaching peak strength. Throughout the heat treatment and machining processes, build orientations were tracked to differentiate between the various build planes (XY, XZ, YZ), tension loading directions (X, Z), and loading-crack growth directions (X-Y, Z-Y) for mode I cracking.

The wrought material was procured as an electric furnace (EF) melted, argon-oxygen decarburized (AOD), and hot-rolled 15.9-mm thick plate produced per ASTM A69344  and supplied in the solution annealed condition. SENT, tension, and excess specimens were excised from the plate and then heat-treated in lab air to either the peak age (482°C for 1 h then air-cooled) or overage (552°C for 4 h then air-cooled) conditions. Tension specimens were oriented with the loading axis parallel to the longitudinal (L) direction (i.e., the plate rolling direction). SENT specimens were oriented with the loading and crack growth in the longitudinal and long transverse directions, respectively; this is the L-T orientation for mode I cracking.46 

Chemical compositions of the as-supplied wrought and as-built (preheat treatment) AM material were measured using inductively coupled plasma optical emission spectroscopy (ICP-OES), glow discharge mass spectroscopy (GDMS), and instrumental gas analysis (IGA), see Table 1. Compositions were nominally similar for the wrought and AM materials. All primary alloy elements, reported in wt%, fell within the ASTM A69344  specifications for wrought 17-4PH. Differences were mainly in the trace elements composition, reported in wppm, with the wrought material having significantly more C, N, V, Mo, Co, and W than the AM material. The AM material overall had smaller concentrations of impurities relative to the wrought material except for a significant amount of O within the powder that decreased slightly during the build. Note that the powder composition was measured by the powder supplier while a commercial laboratory (Evans Analytical Group) measured the wrought and AM build material compositions. The apparent increase in Ni and Si from powder to build is most likely due to lab-to-lab variability.

Table 1.

Elemental Compositions for the Wrought and Additively Manufactured 17-4PH Used in This Study

Elemental Compositions for the Wrought and Additively Manufactured 17-4PH Used in This Study
Elemental Compositions for the Wrought and Additively Manufactured 17-4PH Used in This Study

Mechanical Properties

Hardness and stress/strain properties for the various heat treatments of wrought and AM 17-4PH were reported in Table 2. Prior to hardness testing, samples were incrementally polished with 320, 600, 800, and 1200 grit SiC paper. The Vickers hardness was then measured in accordance with ASTM E38447  using an applied load of 500 g and a dwell time of 10 s. A minimum of nine measurements were taken to establish an average, which was then converted to the Rockwell C hardness scale in accordance with ASTM E140.48  Tension testing was performed on a servohydraulic load frame in accordance with ASTM E8.49  Samples had a 5 mm × 5 mm square gage cross section and were loaded at a constant displacement rate of 8 × 10−4 mm/s, which corresponded to an initial strain rate of ∼1 × 10−4 s–1. The sample elongation was measured using a clip-on extensometer with an 8 mm gage length. Strain hardening parameters (n and α) were determined by fitting the Ramberg-Osgood power law equation50  to the true stress-strain curves.

Table 2.

Mechanical Properties for Wrought and Additively Manufactured 17-4PH

Mechanical Properties for Wrought and Additively Manufactured 17-4PH
Mechanical Properties for Wrought and Additively Manufactured 17-4PH

Microstructure

Specimens for microstructure analysis were cut from excess portions of both the wrought and AM materials. The wrought material was examined in the SA+peak age condition on the LT, LS, and ST material planes, where L, T, and S are the longitudinal, long transverse, and short transverse directions, respectively. The AM material was evaluated on the XY, XZ, and YZ planes for both the HIP+peak age and HIP+SA+peak age conditions. Samples were mounted in cold-curing epoxy-resin and initially ground using 320, 600, 800, and 1200 grit SiC papers, followed by polishing with 3 μm, 1 μm, 0.5 μm, 0.25 μm, and 0.1 μm diamond slurries, and finishing with 0.05 μm alumina or colloidal silica suspension.

Basic grain and defect structure were evaluated using an FEI Quanta 650 scanning electron microscope (SEM). Samples were degaussed prior to imaging and the various forms/planes/heat treatments were evaluated at magnifications ranging from 100× to 5,000×. Secondary electron (SE) imaging was used to identify pores and secondary particles, but the lack of surface deformation and high polish allowed the SE images to also reveal some grain/phase contrast.

Orientation imaging microscopy (OIM) was conducted via EBSD using an FEI Helios G4 UC focused ion beam (FIB)/SEM equipped with an Oxford Instruments Symmetry EBSD detector. OIM maps were collected at two different magnifications using an accelerating voltage of 20 keV and probe current of 3.2 nA. Low-magnification maps were completed at 200× using a 1.0 μm step size, while high-magnification maps were performed at 1,000× using a 0.15 μm step size. Inverse pole figure (IPF) maps and PAG size were determined from the collected OIM data using the MTeX software package.51  The parent austenite grain structure was reconstructed using the approach developed by Nyyssönen, et al.52 

Stress Corrosion Cracking

Fracture mechanics-based testing to characterize the SCC behavior was performed using servo-hydraulic load frames. The crack length was actively monitored using the direct current potential drop (DCPD) method53  and fed into a closed-loop control system to enable stress intensity (K) controlled loading. The SENT specimen geometry, with a 12.20 × 2.67 mm (B/W = 4.6) gage section, was utilized in a pinned-pinned configuration with a linear elastic K calculation.54  Plasticity effects on the mechanical driving force were evaluated55  and determined to be negligible for all evaluated material conditions, justifying the use of K in the current study. Prior to SCC testing, samples were fatigue precracked 0.250 mm out of a 0.940 mm long EDM notch with a root radius of 0.038 mm. Precracks were grown in lab air under cyclic loading at a stress ratio (R) of 0.1 and constant maximum K of .

Environmental control was enabled by isolating the SENT gage section within an acrylic cell capable of flowing liquid solution for immersion testing or gas for atmospheric control. The SENT gage section was polished with 600 grit SiC paper, sonicated in acetone then methanol then deionized (DI) water, and double-coated with a butyl-rubber peel-off lacquer (MICCRO XP-2000) leaving a 5 mm window around the mode I crack path exposed. For inert environment testing, dry nitrogen gas was flowed through the cell at a sufficient rate to maintain a relative humidity (RH) below 5%. During immersion testing, a platinum-coated Nb mesh counter electrode and saturated calomel reference electrode (SCE) were placed in the cell. A neutral, quiescently aerated 0.6 M NaCl solution (pH = 6.2) was flowed from a 2 L reservoir using a peristaltic pump at 20 mL/min and electrochemical potential was applied and/or monitored using a potentiostat operated in floating ground mode. To stabilize the chemistry just prior to SCC testing, the sample was held at OCP for 8 h to 24 h and monitored. The OCP (EOCP) was found to stabilize at −0.190 ± 0.050 VSCE for all examined forms and heat treatments of 17-4PH.

SCC tests were performed at either a fixed (rising K) or a constant K (i.e., ). False crack growth rates caused by DCPD artifacts can occur in positive dK/dt SCC testing,56  and thus tests were performed in both an inert environment to establish false rates/test resolution limits and aggressive environments to capture SCC crack growth rates (da/dt). For the aggressive environment tests, the electrochemical potential (Eapp) of interest was applied for 30 min to 60 min at a K of to establish the target crack-tip environment; testing commenced by increasing K at the aforementioned dK/dt. After rising to the maximum K of interest, the SENT was unloaded at 0.05 mm/min to 0.10 mm/min in displacement control, sonicating in DI water then methanol, heat tinted at 250°C for 1 h to indicate the ending crack length with a straw-brown oxide,57  and fatigued to fracture to avoid ductile fracture distortion at the crack tip. For final K and da/dt calculations, all fracture surfaces were measured to recalibrate the Johnson equation58  crack length calculations with one or two known crack length/DCPD combinations.59  Fractographic features were examined using an SEM operated in SE imaging mode.

Constant K SCC tests were performed at fixed K and several different Eapp; multiple conditions were explored on a single sample and fatigue was performed after each SCC increment to resharpen the crack and enable clear distinction on the fracture surface to confirm cracking. The start of the tests was similar to the rising K tests with precracking, an OCP hold, and a slow rise in K at OCP. Once a maximum K of 30 was reached, the environment was held at an Eapp of interest for a period 8 h to 24 h to capture SCC growth kinetics. At the conclusion of the hold period or after approximately 0.2 mm of crack extension, the environment was returned to OCP and the crack was extended 0.1 mm to 0.3 mm by fatiguing in constant amplitude load control at an initial maximum K of 20 with R = 0.1 or 0.5. This process was repeated a number of times per sample such that bands of fatigue and SCC crack growth were clearly distinguished on the fracture surface. Generally, the applied potential steps started at or near OCP and were driven incrementally more negative (cathodic). Anodic potentials (i.e., Eapp > EOCP) were not evaluated during these constant K SCC tests. To verify critical “takeoff” potentials where SCC susceptibility started, several tests had less cathodic steps repeated to verify shutoff in SCC crack growth. Post-test, the samples were unloaded, cleaned, heat tinted, and fatigued similarly to the rising K experiments. Instantaneous crack growth kinetics were captured by DCPD, but SEM fractography was performed to evaluate and measure the morphologies and actual lengths of the SCC regions. While there was good correspondence between the DCPD and fractography, the reported growth rates were determined by dividing the fractographically measured SCC crack extension by the elapsed time for DCPD-indicated crack extension. The one exception was the AM HIP+peak age alloy, whose fracture surface was sufficiently tortuous that distinguishing the fatigue and SCC regions was not possible, so the DCPD-indicated crack growth rates are reported.

Mechanical Properties and Microstructure

A detailed analysis of the mechanical behavior of the evaluated materials will be presented in a forthcoming publication, but general mechanical properties are provided in Table 2 for the interpretation of current results. As demonstrated in Table 2, both the wrought and AM materials meet the ASTM A69344  mechanical property requirements for the H900 (peak age) and H1025 (overage) tempers. In the peak age conditions, the AM material has higher strength than the wrought alloy, but the strength increase is traded for a reduction in ductility. The aging temperature to reach peak strength is similar in the AM material regardless of whether or not it is solution annealed, but aging with only a HIP treatment results in a significant loss of ductility. Overaging resulted in the nearest match between wrought and AM mechanical properties, but the ductility of the wrought material was still significantly higher than AM. These data demonstrate that the AM material had minimal anisotropy in mechanical properties after heat treatment. Exceptions to this include generally lower ductility in the X direction and slight strength anisotropy in the overage condition. In all cases, 17-4PH undergoes minimal hardening as indicated by the high n in the Ramberg-Osgood fits.

Microstructure evaluation through SEM imaging and EBSD of the wrought and AM materials in the three peak age conditions across the various material planes (i.e., XY, XZ, YZ, LT, LS, and ST) showed no clear indications of microstructural anisotropy. Therefore, only single planes for the AM HIP+peak age, AM HIP+SA+peak age, and wrought SA+peak age conditions are shown in the following microstructure evaluations. Similarly, the overaged microstructures are not shown because the aging process only impacted the nature of the precipitates (not visible at this magnification), thus the grain/martensite structure is effectively identical to the corresponding peak age microstructure. This strong similarity between the overage and peak age condition microstructures at the μm length scale was expected given the relatively stable microstructure of 17-4PH, except for minimal austenite reversion4,60  and precipitate coarsening,61  at the evaluated aging heat-treatment temperatures/times.

To highlight general microstructural differences, OIM maps captured through EBSD are shown in Figure 2. In all conditions, the microstructure was primarily martensitic (colored regions) with negligible austenite phase fractions. At the given length scale, discernable features include the martensite (M) block and PAG structures. In the wrought SA+peak age condition, the M blocks and PAGs were finer than those observed for both the AM HIP+peak age and AM HIP+SA+peak age conditions. With just the HIP treatment, the AM material had the coarsest M block and PAG structure. Adding a SA treatment after the HIP had little effect on the PAG size, but the M block structure appeared to refine to nearly the size of the wrought material.
FIGURE 2.

OIM maps of 17-4PH in the (a,d) HIP+peak age additive, (b,e) HIP+SA+peak age additive, and (c,f) SA+peak age wrought conditions with grain boundaries outlined in black. PAG boundaries are highlighted with the black outlines in (d,e,f).

FIGURE 2.

OIM maps of 17-4PH in the (a,d) HIP+peak age additive, (b,e) HIP+SA+peak age additive, and (c,f) SA+peak age wrought conditions with grain boundaries outlined in black. PAG boundaries are highlighted with the black outlines in (d,e,f).

Close modal
Finer features in the microstructure were revealed by SE imaging at 5,000× magnification, shown in Figure 3. The primary distinguishing feature between the wrought and AM materials at this length scale is the existence of porosity and secondary particles in the AM material vs. the relatively clean microstructure in the wrought material. The AM HIP+peak age material has a combination of porosity and small (<0.25 μm) secondary particles. Some of these particles seem to be aggregated along PAG boundaries, Figure 3(a). Porosity was evident in all evaluated treatments of AM 17-4PH, as indicated by the circular features with light halos in Figures 3(a) and (b). Typical pore sizes were on the order of 0.1 µm to 0.5 μm in diameter and the SA treatment had no discernable effect on the pore structure. In comparison, there are no indications of porosity in the wrought 17-4PH. The small (<0.5 μm), bright, circular features in the wrought material are not pores but more likely Nb carbides known to be present in both wrought62-63  and AM64  17-4PH.
FIGURE 3.

SE micrographs highlighting porosity and secondary particles within additive 17-4PH in the (a) HIP+peak age and (b) HIP+SA+peak age condition vs. (c) wrought in the SA+peak age condition. White arrows indicate particles protruding out of plane.

FIGURE 3.

SE micrographs highlighting porosity and secondary particles within additive 17-4PH in the (a) HIP+peak age and (b) HIP+SA+peak age condition vs. (c) wrought in the SA+peak age condition. White arrows indicate particles protruding out of plane.

Close modal

Stress Corrosion Cracking

Rising K Kinetics

The crack growth rate vs. applied stress intensity relationships measured in dry N2 (RH < 5%) for wrought and AM 17-4PH heat-treated to the peak-aged and overaged conditions are shown in Figures 4(a) and (b), respectively. All tested material conditions exhibit a nominally linear increase in log(da/dt) with increasing K instead of the characteristic rapid acceleration in da/dt typical of stage I SCC crack growth.65  Such behavior has been observed during dry N2 testing in numerous prior studies across a range of alloy systems,56,66-75  which have collectively demonstrated that these linear increases in log(da/dt) with K are not due to real crack extension. Instead, this apparent crack growth is due to geometric contractions of the test specimen76  and localized, crack tip plasticity-induced changes in resistivity73  that act together to increase the DCPD-measured voltage. Prior work56  demonstrated that these “false” crack growth rate vs. stress intensity relationships can be fit to the following equation:
formula
FIGURE 4.

Rising K false crack growth rates for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH at dK/dt = 2 MPa√m/h in an inert dry N2 environment. Growth rates shown are DCPD artifacts associated with crack tip plasticity and specimen compliance not crack extension. Exponential fit lines establish resolution limits for aggressive environment tests.

FIGURE 4.

Rising K false crack growth rates for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH at dK/dt = 2 MPa√m/h in an inert dry N2 environment. Growth rates shown are DCPD artifacts associated with crack tip plasticity and specimen compliance not crack extension. Exponential fit lines establish resolution limits for aggressive environment tests.

Close modal

The fitted C and m coefficient for each condition are shown in Figure 4. Note that X-Y and Z-Y oriented AM specimens for a given heat-treatment condition exhibited similar “false” da/dt vs. K relationships and were therefore simultaneously fit to Equation (1). Nominally identical “false” da/dt responses were observed for wrought and AM 17-4PH heat-treated to the overaged condition, while more significant differences were noted among the tested peak-aged conditions. Specifically, the wrought peak-aged 17-4PH exhibited reduced “false” da/dt for a given applied K relative to the peak-aged AM 17-4PH. The most salient differences were observed between the two peak-aged heat treatments in the AM material, with the “false” da/dt in the HIP+peak age condition exhibiting a ∼4-fold stronger dependence on K (quantified by the m coefficient) than the HIP+SA+peak age condition. The presence of heat treatment-dependent “false” da/dt vs. K behavior is consistent with prior results in other alloy systems. For example, a similar magnitude (∼3-fold) difference in the m coefficient was noted between nonaged and overaged Monel K-500.71 

The finding that the observed linear log(da/dt) vs. K relationships are artifacts and not indicative of real crack growth is corroborated by fractography. Specifically, as shown in Figure 5(a), evidence of crack blunting (indicated by white arrows) was observed on the fracture surface of the HIP+peak age AM material. This observation, coupled with a lack of features suggestive of real SCC extension, confirms the absence of SCC during testing in dry N2 (RH < 5%). Identical blunting features were noted on the fracture surfaces of all other tested material conditions. Interestingly, similar blunting features were also observed across all tested material conditions during slow-rising K experiments conducted in 0.6 M NaCl near OCP (Eapp = –0.190 VSCE), as shown in Figure 5(b) (indicated by white arrows) for the AM HIP+SA+peak age material. Such features (along with similar measured kinetics as the dry N2 results; not shown) demonstrate that 17-4PH under near-OCP conditions in 0.6 M NaCl does not experience SCC crack growth, consistent with expectations from prior literature for 17-4PH immersed in 0.6 M NaCl at ambient temperature.22,77 
FIGURE 5.

Typical morphology of 17-4PH fatigue precrack fronts (identified with white arrows) after rising to K = 30 MPa√m at dK/dt = 2 MPa√m/h with no SCC crack extension. Images shown are for (a) HIP+peak age additive tested in dry N2 and (b) HIP+SA+peak age additive tested in 0.6 M NaCl at Eapp = –0.190 VSCE. Note that the crack extension direction is from left to right and the samples shown were fatigued for further testing after the aforementioned rising K tests.

FIGURE 5.

Typical morphology of 17-4PH fatigue precrack fronts (identified with white arrows) after rising to K = 30 MPa√m at dK/dt = 2 MPa√m/h with no SCC crack extension. Images shown are for (a) HIP+peak age additive tested in dry N2 and (b) HIP+SA+peak age additive tested in 0.6 M NaCl at Eapp = –0.190 VSCE. Note that the crack extension direction is from left to right and the samples shown were fatigued for further testing after the aforementioned rising K tests.

Close modal

A limited number of anodically polarized (i.e., Eapp > −0.190 VSCE) SCC tests were also performed in 0.6 M NaCl. At an applied potential 50 mV anodic of OCP (i.e., −0.140 VSCE), both wrought and AM were not susceptible to SCC as indicated by growth rates (not shown) at or below their respective test resolution limits established in dry N2. At more aggressively anodic potentials, both the wrought and AM materials underwent severe crevice corrosion under the lacquer surrounding the crack path. In multiple cases, this macroscopic dissolution led to failure of the DCPD wires and premature conclusions to the SCC tests. Fractographic inspection of samples showed no clear indication of SCC crack extension.

Cathodic SCC crack growth behavior as a function of stress intensity was captured at Eapp = –1.100 VSCE, shown in Figure 6. Note that false growth rate fits established in Figure 4 were subtracted from the raw data to yield the true crack growth rates following the approach described by Harris, et al.56  This correction also accounted for the effect of deviations from the prescribed loading protocol (discussed below) on the “false” growth rates.74  In general, AM 17-4PH displays higher SCC crack growth rates than wrought 17-4PH across all heat treatments, loading orientations, and applied stress intensities. In the stage II growth region (i.e., nominally K independent growth rate at higher K), the AM material generally cracked 4× to 5× faster than wrought. The only exception is a single test in the AM HIP+SA+Overage (Z-Y) condition where the threshold stress intensity (KTH) was elevated slightly. Noting the generally increased variability in the stage I region (i.e., low K, transient growth), the increased threshold may reflect test-to-test and initial crack tip variability. Another trend is the nearly isotropic behavior in the AM material. The stage II crack growth rates coalesced across the loading orientations (i.e., X-Y and Z-Y) for each of the AM heat treatments. All of the peak-aged AM conditions exhibited similar stage II growth rates, with growth rates only slightly elevated without a solution anneal. The peak-aged threshold stress intensities were consistent across both the wrought and AM materials. However, these thresholds are near the max K () applied to grow the precrack, thus may be influenced by the plastic zone from the precrack. A further minimized precrack plastic zone size (i.e., lower max K) might be required to decouple the apparent Kth,SCC from the precrack. In the overaged condition, growth rates were generally depressed relative to the peak age conditions as expected.20  There is an indication that the overaged AM material may have a slightly higher threshold stress intensity than wrought condition, but the threshold is still low and overshadowed by the generally elevated growth rates in the AM material.
FIGURE 6.

Rising K crack growth rates in 0.6 M NaCl with a cathodic applied potential (Eapp = –1.100 VSCE) for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH at dK/dt = 2 MPa√m/h.

FIGURE 6.

Rising K crack growth rates in 0.6 M NaCl with a cathodic applied potential (Eapp = –1.100 VSCE) for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH at dK/dt = 2 MPa√m/h.

Close modal

Several subtleties associated with the presented cathodic SCC data need to be highlighted. The fracture surfaces in all examined cases were rough and tortuous with increasing amounts of branching (as indicated by out-of-plane features) with higher applied K. The ending crack lengths were optically measured post-test and consistently found to be on the order of 1.5× longer than DCPD-indicated final crack lengths. This underprediction by in situ DCPD necessitated the previously described two-point, post-test corrections to the crack length and associated K, da/dt calculations. The fatigue precrack front, discernable by fatigue-SCC morphology change, and ending SCC crack length, marked by heat tinting, were used as reference points in this analysis. Functionally, this means that the reported is only strictly accurate for the low K results. Specifically, as the K increases, the dK/dt increased up to a maximum of . However, this variable dK/dt had little impact on growth rates as indicated by the less than 2-fold and 3-fold change in growth rate beyond for the peak-aged and overaged materials, respectively.

Constant K Kinetics

The effects of various applied potentials were examined through a series of constant K SCC tests in 0.6 M NaCl, see Figure 7. Based on the rising K tests, a K of to was chosen to target the stage II crack growth region. For this testing, an absence of measurable crack growth after an 8 h to 24 h test period for a given Eapp was used to identify nonsusceptible region, indicated by a crack growth rate of 1 × 10−7 mm/s and a downward arrow in Figure 7. The crack growth kinetics observed during constant K testing can be dependent on load and polarization history (due to blunting, charging, or transient behavior), as such all reported data in Figure 7 were measured in the “freshly fatigued” condition (i.e., polarized within 30 min of fatigue).
FIGURE 7.

Constant K SCC crack growth rates as a function of applied potential in 0.6 M NaCl for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH. Where no crack extension was observed, a growth rate of 1 × 10−7 mm/s with a down arrow is shown.

FIGURE 7.

Constant K SCC crack growth rates as a function of applied potential in 0.6 M NaCl for various forms and loading orientations of (a) peak-aged and (b) overaged 17-4PH. Where no crack extension was observed, a growth rate of 1 × 10−7 mm/s with a down arrow is shown.

Close modal

The general applied potential dependence of the SCC susceptibility in Figure 7 is consistent with prior reports on the SCC behavior of high-strength steels.67-68  Specifically, high resistance to SCC is observed at mildly cathodic potentials that then transitions to significant SCC susceptibility as the applied potentials become more cathodic. Critically, the border between these nonsusceptible and SCC-susceptible regions provides a “take-off” potential that can be used as a proxy to rank the relative SCC susceptibility of each form/heat treatment of 17-4PH. Stark differences in this take-off potential metric were found for the wrought and AM materials in the peak-aged conditions. As shown in Figure 7, AM 17-4PH exhibited take-off potentials that were between 200 mV and 400 mV more positive than the wrought condition. The peak-aged AM condition that did not undergo a SA step was the most susceptible, with an estimated takeoff potential between −0.400 VSCE and −0.500 VSCE. However, with a solution anneal, the peak-aged AM takeoff potential shifted to a slightly more cathodic potential (–0.600 VSCE), but still exhibits rapid growth rates (∼1 × 10−3 mm/s) in the −0.600 VSCE to −0.800 VSCE range. Interestingly, the peak-aged AM condition has a distinct transition at cathodic potentials beyond −0.800 VSCE. The crack growth rate decreases by an order of magnitude and then follow the general trend for the peak-aged wrought condition, though at an elevated rate. Multiple loading orientations were evaluated for the AM material and the SCC behavior was isotropic across all applied potentials and heat-treatment conditions.

In the overaged condition, the AM 17-4PH exhibits a slightly less cathodic takeoff potential than wrought (i.e., indicating increased SCC susceptibility) and generally higher crack growth rates in the susceptible range, though the magnitude of the differences are less pronounced than was observed for the peak-aged condition. Comparison of the peak-aged and overaged AM conditions shows that overaging shifts the takeoff potential more negative, but similar rapid crack growth rates are observed at highly susceptible conditions (e.g., –1.200 VSCE to −1.400 VSCE). Conversely for wrought 17-4PH, overaging has a dual benefit of improving the takeoff potential and simultaneously depressing the crack growth rate for a given applied potential.

Fractography

Fractography showed that differing loading rates, orientations, and K values did not impact the fracture morphology, except for the observation of increased branching and tortuosity at high K. Differences in morphology were primarily observed as a function of form, heat treatment, and applied potential, which are detailed below and summarized in Table 3. Regarding the peak-aged condition, representative fractographs are shown in Figure 8. For peak-aged wrought 17-4PH, crack growth was almost exclusively transgranular (TG). Though only shown at Eapp = –1.100 VSCE, Figure 8(a) is representative of the fracture morphologies observed for peak-aged wrought 17-4 across all susceptible Eapp. Conversely, peak-aged AM 17-4PH generally exhibited mixed IG+TG failure, with the fraction of IG depending on the applied potential and SA protocol. For example, between the take-off potential and −0.800 VSCE, a significant fraction of IG was observed for both HIP+SA+peak age and HIP+peak age, with the former exhibiting mostly IG fracture (Figure 8[d]). However, as the applied potential became more cathodic, the extent of TG fracture increased for both conditions, with HIP+peak age (Figure 8[b]) exhibiting mixed TG and IG fracture and HIP+SA+peak age (Figure 8[c]) transition to largely TG fracture. Interestingly, the Eapp where the AM HIP+SA+peak age condition began to exhibit increasingly TG-dominant crack growth (≈ −0.925 VSCE) closely corresponds to the unexpected reduction in growth rate shown in Figure 7(a).

Table 3.

SCC Morphologies Evaluated as a Function of Applied Potential and 17-4PH Form/Heat Treatment

SCC Morphologies Evaluated as a Function of Applied Potential and 17-4PH Form/Heat Treatment
SCC Morphologies Evaluated as a Function of Applied Potential and 17-4PH Form/Heat Treatment
FIGURE 8.

SCC morphologies for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the (a) wrought SA+peak age, (b) additive HIP+peak age, and (c) additive HIP+SA+peak age conditions. At Eapp = –0.800 VSCE, additive in the HIP+SA+peak age condition exhibited the morphology shown in (d).

FIGURE 8.

SCC morphologies for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the (a) wrought SA+peak age, (b) additive HIP+peak age, and (c) additive HIP+SA+peak age conditions. At Eapp = –0.800 VSCE, additive in the HIP+SA+peak age condition exhibited the morphology shown in (d).

Close modal
In the overaged condition, the wrought and AM materials both exhibit primarily IG crack growth across all applied potentials, shown in Figure 9. For wrought 17-4PH, IG was mixed with nearly evenly spaced small regions of TG fracture. The same was true for the AM material, but the TG regions were further spaced, resulting in IG being more dominant. Interestingly, the overaged wrought material exhibited a number of larger-scale out-of-plane fissures parallel to the LT plane at highly cathodic potentials (image not shown); these features were not observed in the AM material. Given the plane and spacing, these fissures might have been following delta ferrite stringers formed during hot-rolling.78-79  While specific characterization aimed at identifying delta ferrite stringers was not performed for the current study, stringers were noted in the wrought plate during other characterization activities (and not in the AM material).
FIGURE 9.

SCC morphologies for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the (a) wrought SA+overage and (b) additive HIP+SA+overage conditions.

FIGURE 9.

SCC morphologies for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the (a) wrought SA+overage and (b) additive HIP+SA+overage conditions.

Close modal
High-magnification micrographs of the intergranular facets observed for the various tested conditions are shown in Figure 10. As previously mentioned, the wrought SA+peak age condition was unique in that intergranular fracture was not broadly observed. Conversely, both AM HIP+SA+peak age and AM HIP+peak age exhibited significant IG fracture between the take-off potential and −1.000 VSCE. With a HIP+SA treatment (Figure 10[a]), the peak-aged AM material had IG facets decorated with 0.1 µm to 0.5 μm pore-like features and sub-surface particles less than 100 nm in size. The remainder of the facet was relatively smooth. Without the SA treatment, the peak-aged AM material had a combination of the pore-like features and particles protruding from the facets (Figure 10[b]). The small, nm-scale particles appear to be increased in size and number relative to the solutionized AM material. Additionally, protruding crater-like features that surround the particles and instances of broken particles may indicate that they may play a role in the fracture of the facet. In the overage conditions (Figures 10[c] and [d]), fracture was dominated by IG for both the AM and wrought material. Unlike the peak-aged materials, the facets are less smooth and are covered with features that suggest local plasticity. The porosity is still evident on the AM facets while the wrought facets exhibit no evidence of porosity.
FIGURE 10.

Intergranular SCC facets for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the additive (a) HIP+SA+peak age, (b) HIP+peak age, (c) HIP+SA+overage conditions, and (d) wrought in the SA+overage condition.

FIGURE 10.

Intergranular SCC facets for 17-4PH cathodically polarized in 0.6 M NaCl at Eapp = –1.100 VSCE in the additive (a) HIP+SA+peak age, (b) HIP+peak age, (c) HIP+SA+overage conditions, and (d) wrought in the SA+overage condition.

Close modal

The preceding results demonstrate that AM 17-4PH exhibits increased SCC susceptibility relative to wrought 17-4PH at comparable yield strength levels across all evaluated heat treatments, environments, and loading schemes. Clear differences in both metallurgical features and fracture morphology were also noted across the tested conditions. Three questions arise from these results:

  1. What is responsible for the observed increase in SCC susceptibility of AM 17-4PH?

  2. What is responsible for the observed changes in fracture morphology as a function of applied potential?

  3. What is responsible for the observed more positive take-off applied potential for the HIP+peak age relative to the HIP+SA+peak age AM conditions?

In the following discussion, a decohesion-based failure framework for hydrogen-driven SCC is used to link the observed susceptibility trends to changes in the governing microstructure features that are induced by the AM process. Based on this mechanistic framework, the contribution of the observed differences in PAG size, hydrogen-metal interactions, compositional heterogeneities, second-phase particles, and porosity to the SCC susceptibility of the wrought and AM materials are considered. The implications of these evaluations are then discussed in the context of possible process modifications that may improve the SCC resistance of AM 17-4PH.

Decohesion-Based Mechanistic Framework for Stress Corrosion Cracking

It is well-established that hydrogen embrittlement is responsible for SCC in high strength steels exposed to marine environments under cathodic polarization.67-68,80  A litany of theories have been proposed to describe the microscale mechanism(s) responsible for such hydrogen-induced degradations in fracture resistance.81-84  Recent literature strongly suggests that hydrogen-assisted fracture occurs via interface decohesion that is synergistically driven by the combined interactions of hydrogen-induced reductions in cohesive strength and hydrogen-modified deformation behavior.81-82,85  While it has known shortcomings that preclude quantitative prediction, the decohesion-based model for KTH proposed by Gerberich83  (and subsequently modified by others69-70 ) phenomenologically captures the general dependencies of SCC on hydrogen and microstructure-based parameters.66,73,86  As such, this framework can be leveraged to understand the mechanistic influence of possible AM-induced microstructural changes:
formula
where kIG is the Griffith toughness of the susceptible interface (e.g., martensite lath interfaces, grain boundaries, etc.) which can be modified (either negatively or positively) by the presence of trace impurities. The influence of these impurities is described by , where Ci and βi are the local concentration and potency of the effect of the ith impurity on the intrinsic Griffith toughness. α describes the degradative potency of hydrogen, while C is the hydrostatic stress-enhanced hydrogen concentration that is pertinent to the fracture process zone. Last, σys is the alloy yield strength and the α″ and β′ parameters are coefficients determined from modeling of dislocation shielding on the crack tip elastic stress.

It is important to recognize that the variations in SCC susceptibility observed in Figures 6 and 7 are represented by differences in crack growth kinetics while Equation (2) describes the dependence of KTH on hydrogen content and microstructure-based parameters. However, the authors propose that the Gerberich model for KTH can be more generally interpreted, and thus extended to serve as the description of the conditions required for hydrogen-assisted crack advance or a proxy for interface failure stress (i.e., a failure criteria). This paradigm allows Equation (2) to then serve as a platform to qualitatively assess the observed variation in SCC kinetics in the context of potential AM-induced microstructure modifications. The extension of this general framework to understand how AM-specific microstructure features may impact the susceptibility of SCC crack growth is justified by three considerations. First, recent work has extended the Ritchie-Knott-Rice (RKR) paradigm for crack advance87  to hydrogen environment-assisted cracking (HEAC) conditions.56  The RKR model proposes that crack advance occurs when a local failure criterion (e.g., critical stress87  or strain88 ) is exceeded over a critical distance (xcrit) ahead of the crack tip. In a HEAC paradigm, this failure criterion is degraded by the local hydrogen content.56,89  The authors propose that the Gerberich relationship (Equation [2]) describing the dependence of the continuum-scale threshold (e.g., KTH) on microstructure-dependent parameters is analogous to, and thus can be extended to broadly capture, how the local failure criteria at xcrit is impacted by microstructure variations. Second, it is recognized that the salient role of H diffusivity for HEAC is not captured by the Gerberich framework. However, (1) this omission will be explicitly considered by discussion of the AM-specific H-material interactions that will impact C, and (2) prior work has demonstrated synergy between micromechanical models that describe the microstructure dependence of KTH and crack growth kinetics.67,69,90-92  Finally, this framework will not be used for quantitative predictions; as such, the rigor of the functional form of Equation (2) will not compromise the use of this framework to postulate how microstructure variations may impact the failure criteria.

Once justified, evaluation of the parameters in Equation (2) highlights several factors that may influence the relative SCC susceptibility of AM vs. wrought 17-4PH. First, differences in yield strength could have an impact, with higher yield strength leading to a reduced failure criterion. Second, a difference in hydrogen-material interactions, such as the increase in the diffusible hydrogen concentration at a given applied electrochemical potential, could increase the hydrogen available to participate in the fracture process, leading to an increase in C and a decrease in the failure criterion. Third, differences in the near-crack deformation behavior or slip morphology that impede dislocation emission would lower the failure criterion through modifications in the α″ and β′ terms. Fourth, variations in the trace impurity content at the susceptible interface, as would be expected for both differences in aging condition and AM/wrought stock, could either degrade or increase kIG via modification of the term in Equation (2), leading to a concomitant decrease or increase the failure criterion. Finally, variations in the local stress state in the fracture process zone at the crack tip (particularly modifications in the hydrostatic stress distribution as may be induced by porosity, inclusions, etc.), could serve to locally affect C and therefore modify the failure criterion. It should be noted that the AM and wrought materials tested in the current study were specifically heat-treated to minimize variations in the yield strength. For example, as shown in Table 3 and Figure 7, the AM HIP+SA+peak age material has a ∼100 MPa increase in yield strength relative to the AM HIP+peak age material (which is nominally identical to the wrought SA+peak age material), but a −200 mVSCE more negative take-off potential than the AM HIP+peak age alloy. As such, yield strength differences are unlikely to be the source of the observed variations in SCC susceptibility and the following discussion in the Microstructural Origins of Degraded Stress Corrosion Cracking Performance section will focus on the other identified parameters.

In addition to understanding the contribution of microstructural features, Equation (2) can also be leveraged to mechanistically understand the observed changes in fracture morphology as a function of applied potential for a given alloy condition (Table 3). First, kIG in Equation (2) is not limited to describing the intrinsic Griffith toughness of a grain boundary; it can be used to describe the intrinsic toughness of any crack pathway.93  Second, it is unlikely that the αC term in Equation (2) will be constant for all possible crack pathways under a constant environmental condition. For example, the literature demonstrates that the driving force for hydrogen segregation to Σ3 boundaries is low,94-95  which would act to strongly reduce the hydrogen present at this interface (and therefore reduce C). While the specific details of the hydrogen trap binding energies for the PAG boundaries and the martensite lath interfaces in the current materials are not known, the literature suggests it is reasonable to expect variations in the extent of hydrogen trapping between these features.96  Similarly, due to the intrinsic differences between these interface features (e.g., PAG boundaries and martensite laths), it is also likely that the degradative effects of hydrogen (i.e., α in Equation [2]) will be distinct for each.

Based on the above discussion, it is likely that kIG and αC vary between possible fracture pathways, which would lead to different fracture paths becoming favorable at different applied potentials. This variation in fracture pathway was specifically noted for the AM HIP+SA+peak age condition (Table 3), where the fracture morphology was observed to transition from IG to TG-dominant fracture as the applied potential became increasingly negative. Considering Equation (2), this could manifest if PAG boundaries in the AM HIP+SA+peak age material are described by a low kIG and a modestly applied potential dependent αC term, but the TG fracture pathway is characterized by a high kIG and stronger applied potential dependence of αC. Such a scenario is illustrated schematically in Figure 11 for two conditions of peak-aged 17-PH.
FIGURE 11.

(a) Schematic illustrations of the primary grain boundary interfaces in 17-4PH and (b) the postulated response of those interfaces’ failure stresses to stress-assisted hydrogen concentration (C) in peak-aged 17-4PH. The IG and TG fractographic bases for these postulations are superimposed for reference.

FIGURE 11.

(a) Schematic illustrations of the primary grain boundary interfaces in 17-4PH and (b) the postulated response of those interfaces’ failure stresses to stress-assisted hydrogen concentration (C) in peak-aged 17-4PH. The IG and TG fractographic bases for these postulations are superimposed for reference.

Close modal

For AM in the HIP+SA+peak age condition, the failure stress required for interface decohesion at low hydrogen concentrations (which trends to kIG) must be higher for the martensite lath boundaries than the PAG boundaries given the observed IG fracture at more positive applied potentials. As the hydrogen content is increased, the interface failure stress is reduced for both boundaries, but at different rates (e.g., due to different α values), resulting in a cross-over. This postulated cross-over potential is likely associated with the observed unexpected decline in crack growth kinetics for the AM HIP+SA+peak age material at −0.950 VSCE (Figure 7[a]). As shown in Table 3 and Figure 8, the predominant fracture mode for the AM HIP+SA+peak age material changes from IG to predominantly TG between −0.800 VSCE and −1.100 VSCE. It is therefore reasonable to suggest that there is an intermediate potential between these two bounding conditions where both IG and TG fracture pathways are (effectively) equally favorable. If multiple fracture pathways are equally favorable, this will likely lead to an increasingly tortuous, mixed fracture pathway that reduces the local driving force for crack advance (i.e., due to crack branching and/or deflections) and therefore causes a decrease in the crack growth rate. Indeed, a mixture of the fracture morphologies shown in Figures 8(c) and (d) was observed at −0.950 VSCE (image not shown). However, once the applied potential became sufficiently negative that a single fracture pathway (a different fracture path in this case, Figure 8[c]) was dominant, then the crack growth rate increased from the locally reduced value, as observed in Figure 7(a).

Conversely, wrought in the SA+peak age condition exhibited a consistent TG fracture pathway not subject to the aforementioned variations seen in AM 17-4PH (Table 3). Accordingly, the wrought growth rate steadily increased as the applied potential became increasingly negative (Figure 7[a]). These observations might point to a higher PAG kIG relative to martensite lathe boundaries in wrought as compared to AM and/or an attenuated H degradation of wrought PAGs as described by α. Clearly, other factors are likely to contribute to these processes (e.g., differences in trace impurity contents, local deformation, etc.) and high-fidelity atomistic simulations are required to confirm the above postulations. However, this thought experiment highlights the utility of the current decohesion-based mechanistic framework (Equation [2]) to postulate mechanistic underpinnings for observed behavior.

Microstructural Origins of Degraded Stress Corrosion Cracking Performance

Grain Boundary Impurity Segregation

The current investigation shows that IG fracture in AM 17-4PH is associated with both enhanced SCC susceptibility and accelerated crack growth rates relative to the wrought material. In the peak age condition, this phenomenon is especially evident as wrought 17-4PH exhibits a transgranular fracture morphology (Table 3), whereas the AM material had increased SCC susceptibility that correlated with IG fracture (Table 3 and Figure 7[a]). Comparison of the various fracture morphologies (Figures 8 and 9) and grain structures (Figure 2) shows that the intergranular fracture is occurring along the PAG boundaries. Generally, fracture along PAGs in steels during ambient temperature deformation can be promoted by (1) elemental segregation of deleterious impurities during thermal treatment, (2) nucleation of secondary phases on the grain boundary, and/or (3) aggressive environments.97  The latter was directly evaluated through varying Eapp (Figure 7), while the former two can be grouped under a phenomenon called temper embrittlement.98 

Under certain thermal processing schedules, many steels exhibit a transition from ductile microvoid coalescence or TG fracture to IG-dominated brittle fracture. For example, this “temper embrittlement” has been reported in wrought 17-4PH after prolonged exposure (on order of months) to temperatures between 300°C and 450°C.99-101  It is notable that this embrittlement occurs in the absence of an aggressive environment (though hydrogen-rich environments can exacerbate the effect102 ) and is largely attributed to the segregation of impurity elements to PAG boundaries, including: Si, Ge, Sn, N, P, As, Sb, Bi, S, Se, and Te.97  Mechanistically, these segregated impurities act to degrade the grain boundary toughness (i.e., kIG in Equation [2] with the impact of the impurities captured by the term), leading to reduced fracture resistance. Critically, this embrittlement has also been observed under more conventional heat-treatment schedules (e.g., the H1025 temper) for 17-4PH with phosphorus contents of ∼0.03 wt%.103 

Considering the current alloys, the postbuild AM material does have 6-fold more S and 65% more Si than the wrought material (Table 1), though these levels are still well within the acceptable limits for 17-4PH per ASTM A693.44  However, given the differences in thermal processing between the AM and wrought alloys, it is possible that even larger differences in concentration are present at grain boundaries. For example, >1,000-fold increases in the grain boundary S concentration over the bulk content were recently measured in Monel K-500 using Auger electron spectroscopy.71  This local agglomeration of impurities at the grain boundary can be exacerbated by three factors: (1) large grain size,104  (2) long-term exposures to intermediate temperatures,105  or (3) slow cooling rates after high-temperature heat treatments.106  Regarding grain size effects, because larger grains result in a lower grain boundary area, such grains will likely have an increased grain boundary impurity content than small grains for a constant bulk impurity content. Conversely, increased time at intermediate temperatures and/or slow cooling rates allow for the diffusion of impurities toward grain boundaries, increasing the grain boundary impurity content.

Compared to the wrought alloy, the AM material in this study has several features that likely promote elemental segregation including a larger PAG size (Figure 2) and the slow 310°C/h cooling rate from the HIP treatment. For example, the HIP+peak age condition, which had the worst SCC susceptibility, has both of these contributing factors. Conversely, the wrought peak-aged material has the smallest PAG size (Figure 2), did not undergo the HIP treatment, and has improved SCC resistance. The improved SCC resistance of the AM HIP+SA+peak age condition, that underwent a high-temperature (1,040°C) SA treatment followed by air cooling, exhibited improved SCC resistance relative to the AM HIP+peak age alloy (Figure 7[a]). This is also consistent with a likely role of temper embrittlement. Specifically, the SA treatment will act to both rehomogenize grain boundary impurities and refine the PAG size (as demonstrated in Figure 2), while the air-cool would reduce the extent of resegregation back to the grain boundaries during cooling (as it is significantly faster than the 310°C/h cooling rate used in the HIP treatment). Similar trends in behavior are observed for the overaged alloys, with the AM HIP+SA+overage condition exhibiting higher SCC susceptibility than the wrought SA+overage condition (Figure 7[b]). Last, the increased observation of IG fracture in both the AM and wrought overaged material relative to the peak-aged alloys is also consistent with temper embrittlement effects. Tempering to the overaged condition involves holding the specimen at 552°C for 4 h, while tempering to the peak-aged condition requires 1 h at 460°C/482°C. Such prolonged exposure would exacerbate impurity segregation in the overaged condition, increasing the likelihood of IG fracture pathways being favorable relative to TG pathways.

Porosity

Review of the microstructure (Figure 3) and fractography (Figure 10) shows that sub-μm pores are omnipresent across all evaluated AM conditions. It is likely that this porosity is due to insoluble Ar107  that was introduced both during powder atomization (as the atomizing gas) and during the build itself (as the inert build atmosphere). Porosity is a well-established challenge for AM alloys35-36,108-111  and is associated with a general reduction in toughness and ductility.112  However, it should be noted that the pore size range of interest in AM materials is typically 5 µm to 20 μm.33,35-36,113  The 0.1 µm to 0.5 μm pores observed in this study (Figures 3[a] and [b]) are well below this typically evaluated range and, more critically, appear to be resistant to HIP treatments. As such, it is likely that such sub-micrometer pores will be present in AM 17-4PH components. However, understanding of how these 0.1 µm to 0.5 μm pores affect crack propagation behavior is highly limited, especially with regard to SCC.

It is logical to suggest a deleterious influence of this sub-μm porosity on the SCC resistance given the large number of pores present on the fracture surfaces of all tested AM conditions (both TG and IG; Figure 10). Interestingly, qualitative examination of the voids observed on the fracture surface indicates little change in the void shape and size relative to the pores noted in the metallurgical cross sections (Figure 3). This qualitative similarity suggests minimal damage evolution occurred proximate to the voids during crack advance, consistent with the theoretical expectation that small voids will grow at slower rates relative to large voids.114  Critically, such minimal damage evolution suggests a brittle fracture process that, coupled with the discontinuous nature of subcritical SCC growth, provides a plausible mechanistic pathway by which these voids may contribute to the generally increased SCC susceptibility of AM 17-4PH (Figure 7). Specifically, the interaction of these pores with hydrogen in the fracture process zone likely impacts how these voids contribute to SCC. Hydrogen is expected to agglomerate proximate to/within the pores in the fracture process zone due to both (1) the local enhancement in hydrostatic stress increasing local solubility115  and (2) hydrogen trapping at the pores. At low concentrations, hydrogen has been shown to accelerate void growth and nucleation processes,116  but as the hydrogen concentration is increased to elevated levels (as would be expected proximate to the crack tip115 ), a transition to brittle fracture modes occurs.85,117  As such, it is possible that this local elevation in hydrogen content could reduce the intrinsic fracture resistance that, coupled with the elevated local stresses proximate to the voids (because they are stress concentration sites), may result in easier crack initiation at pores.

The observed increases in crack growth rate under fixed K loading for all AM conditions relative to their wrought counterparts (Figure 7) directionally support this proposed role of porosity. This can be rationalized using the aforementioned RKR model for crack propagation,87  where the fracture resistance (which can be degraded by hydrogen)89  must be exceeded over a critical distance in order for crack extension to take place. If the proposed hydrogen-induced crack initiation occurs at these voids, the high number density of such features (Figures 3 and 10) indicates that it is plausible a void will reside between a given crack tip position and the critical distance associated with crack advance. These preinitiated cracks could act to reduce the critical length (xcrit) that must be exceeded for crack propagation or vary the critical stress/strain needed for failure, leading to a general increase in both SCC susceptibility and the crack propagation rate, consistent with the up-shift in da/dt and more positive “take-off” potentials noted in Figure 7. This proposed contribution of porosity is also directionally consistent with previously invoked mechanisms for hydrogen-induced fracture in TWIP steels118  and IN718.119  Specifically, in these materials, it is postulated that voids nucleated at intersecting slip bands or deformation twin/grain boundary junctions serve as crack initiation sites that facilitate subsequent crack propagation. While direct evaluation of this proposed mechanism is outside the scope of the current study, future efforts will focus on confirming this mechanistic role of voids via high-fidelity microstructure characterization and targeted modeling approaches.

Secondary Phase Particles

In addition to porosity, examination of the fractographs (Figure 10) and metallurgical cross-section micrographs (Figure 3) also reveals the presence of sub-μm size secondary phase particles. Literature establishes that the most commonly observed particles that form in conventional 17-4PH are NbC and Cr23C6,62-63,100,120  while these phases along with Cr7C3,121  MnS,122  and oxide inclusions32-33,123  have been reported in AM 17-4PH. Of these, NbC, Cr23C6, and oxide inclusions are the most likely secondary phase particles in the current study. Regarding NbC, Nb is a strong carbide former that is generally added to sequester C to reduce Cr23C6 nucleation.64,124  NbC particles are observed after all standard heat treatments for 17-4PH,63-64,120  consistent with its reported solvus of >1,250°C.64  However, it is generally considered to be nondetrimental from a mechanical property perspective due to its low volume fraction and large nearest neighbor distances.

Conversely, grain boundary precipitation of Cr23C6 is associated with increased corrosion susceptibility in stainless steels due to the local depletion of the Cr content.125  Critically, while Nb is added to suppress Cr23C6 formation, the Nb content typical of 17-4PH (<0.3 wt%) is not high enough to completely prevent nucleation of this deleterious phase during common heat treatments.63-64  One challenge associated with Cr23C6 is that volume fractions are generally low, making detection difficult via standard characterization techniques like x-ray diffraction.64  As such, it is most common for this phase to be reported after long-term (months) of long-term exposures at 300°C to 400°C (meant to replicate in-service conditions at nuclear powerplants),100  unless specific characterization approaches meant to highlight this secondary phase are performed (e.g., double-loop electrochemical potentiodynamic reactivation tests [DL-EPR]).63  For example, these DL-EPR experiments have demonstrated the presence of Cr23C6 after aging at or above 500°C in wrought 17-4PH.63  Phase diagrams for 17-4PH64  predict that Cr23C6 will start precipitating at ∼700°C for the Nb content in the current alloys; these results suggest that the pertinent temperature range for Cr23C6 formation in 17-4PH under typical heat treatments is 500°C to 700°C.

Regarding the currently used high-temperature heat treatments, the HIP treatment was conducted at 1,125°C, which is above the Cr23C6 solvus.64  However, given the slow cooling rates associated with the HIP treatment (310°C/h), an appreciable amount of time was spent in the noted temperature range where Cr23C6 precipitation is favorable. Conversely, the SA treatment (conducted at 1,040°C) is also above the Cr23C6 solvus,64  but was followed by air cooling, which would drastically reduce the time spent in the temperature range where Cr23C6 precipitation is favorable. Considering these differences in the context of the SCC susceptibility noted for the peak-aged AM materials, it is interesting to note that the HIP+SA+peak age exhibited increased SCC resistance relative to the HIP+peak age condition. This improvement, coupled with the notable reduction in secondary phase particles on the fracture surface for the HIP+SA+peak age relative to the HIP+peak age condition (Figure 10), suggests that these Cr23C6 carbides may be contributing to the increased SCC susceptibility of the AM materials. Mechanistically, it seems reasonable that a process similar to that previously postulated for the voids observed in the AM material would be operative for the Cr23C6 particles. Specifically, the presence of hydrogen and the increased local stress proximate to these particles would act to make crack initiation more favorable, leading to enhanced SCC susceptibility. Such a scenario has been explicitly invoked in prior modeling of hydrogen-assisted cracking of carbide-containing steels.126 

As previously discussed, the SA treatment should have attenuated elemental segregation and Cr carbides precipitated during the HIP, yet the SCC growth kinetics of the AM material in the HIP+SA+peak age condition still underperform wrought (Figure 7[a]), suggesting a role of other factors. While porosity appears to be the most likely contributor, it is pertinent to note that several LPBF stainless steel studies have shown that oxygen introduced through the build atmosphere or powder can result in finely dispersed Mn/Si/Nb/Cu-rich oxides on the order of 0.05 µm to 0.10 μm in size.30-33  These oxides can occur with as little as 260 wppm of dissolved O,31  and have been theorized to be electrochemically active during SCC, albeit under supercritical water conditions.30  The current lot of AM material has multiple of these hallmarks including a 600 wppm/480 wppm O concentration in the powder/build (Table 1) and second phase particles less than 0.25 μm in size (Figure 3[a]) that are insoluble up to 1,040°C (Figure 10[a]). It is expected that these insoluble second phase particles are oxide inclusions that, in a similar manner to Cr23C6 carbides, enable easier crack extension along the grain boundary in AM 17-4PH. While literature indicates that the electrochemical activity of these oxides is likely minimal under the dilute neutral chloride solution conditions,127  it is possible that the observed easier crack extension observed in the AM HIP+SA+peak age condition may be facilitated by an influence of these fine oxide particles on the local crack tip stress field. However, like the aforementioned discussion on sub-μm porosity, additional studies are needed to understand the mechanistic details regarding how the oxide inclusions and other secondary phase particles may affect the local crack-tip fields.

Hydrogen-Metal Interactions

In addition to the above microstructural features, it is also likely that differences in hydrogen-metal interactions such as diffusivity, solubility, and trapping may exist between wrought and AM 17-4PH. While a detailed study of these parameters is ongoing for these alloys, it is still useful to comment on (1) what specific parameters would be mechanistically impactful and (2) how such impacts would arise. First, it is well established across multiple alloy systems that SCC susceptibility is directly tied to the hydrogen concentration available to participate in the fracture process (i.e., the diffusible hydrogen concentration).115  As such, the observed differences in cracking kinetics could be related to differences in the diffusible hydrogen concentration (captured by C in Equation [2]) for a given electrochemical potential between the AM and wrought conditions. Figure 3 demonstrates that subtle, but tangible differences in microstructure exist between AM and wrought 17-4PH. Such subtle microstructural variations can induce significant differences in the diffusible hydrogen concentration, as demonstrated in a prior study that assessed the hydrogen uptake behavior of five separate heats of a Ni alloy that were all heat-treated in accordance with the governing procurement specification.128  Regarding AM materials, a recent study by Bertsch, et al., has evaluated the hydrogen uptake behavior of austenitic stainless steels manufactured using LPBF, direct energy deposition (DED), and conventional processes.129  Critically, these authors report a systematic increase in hydrogen uptake for both AM (LPBF and DED) materials relative to the conventional alloy. If such an increase were to be observed for the 17-4PH, this would potentially explain the enhanced SCC susceptibility of all tested AM 17-4PH materials relative to their wrought counterparts. Moreover, Bertsch, et al., also reported tangible differences in hydrogen uptake between the two AM materials in their study, consistent with the variations observed for different processing approaches (albeit postbuild) in Figure 7. Studies are now underway to measure the diffusible hydrogen concentration as a function of applied potential in the current materials and will be reported in a forthcoming manuscript.

Second, it is also possible that differences in hydrogen diffusivity could contribute to the differences in crack growth rates shown in Figure 7. The previously highlighted RKR failure criterion for crack advance is independent of the diffusivity, but the rate at which the criterion is met (i.e., the rate of crack advance) is explicitly linked to diffusivity. Given the accelerated crack growth kinetics for the AM material across all tested conditions, this would suggest that the AM material exhibits an increased hydrogen diffusivity relative to the wrought condition. Prior studies comparing conventional and AM austenitic stainless steels have reported both increased130-131  and reduced132  hydrogen diffusivities in AM materials. Interestingly, these studies claimed an important role of the subgrain boundaries introduced by the AM processing in either accelerating131  or slowing132  hydrogen diffusion. While measurements of the hydrogen diffusivity in AM vs. wrought 17-4PH are not available in the open literature, these studies in austenitic steels highlight the potential role of AM-induced microstructural features in modifying diffusion behavior. Considering the current study, hydrogen trapping at the widespread sub-μm porosity observed in the AM alloys (Figures 3 and 10) would be reasonably expected to impact the hydrogen diffusivity. However, such traps would be expected to slow hydrogen diffusion,115  which would result in reduced crack growth kinetics. This expectation is directionally opposed to the observed increase in crack growth rate for all tested AM 17-4PH materials (Figure 7) as a function of applied potential, underscoring the need for direct measurements of the hydrogen diffusion and trapping behavior in these alloys.

Crack-Tip Deformation Behavior

While the above “bulk” characteristics will play an important role in establishing SCC susceptibility, SCC ultimately proceeds via highly localized interactions between hydrogen, deformation, and microstructural features within the fracture process zone proximate to the crack tip.115  As such, local characterization of the near-crack region is necessary to determine differences in near-crack deformation as a function of electrochemical potential (i.e.,  hydrogen content) and heat-treatment condition. Critically, any observed variations in the local deformation behavior can then be leveraged to understand the microscale mechanisms governing SCC susceptibility. The influence of this local deformation is phenomenologically captured in Equation (2) by the α″ and β′ terms and mechanistically contributes to crack advance under the RKR paradigm via effects on the local mechanical driving force. Prior studies of the near-crack region suggest that hydrogen exposure results in a more localized deformation distribution proximate to the crack for a given driving force.81,133-134  However, considering the current alloys, how this hydrogen-enabled localization is affected by the widespread sub-μm porosity, oxide inclusions, and carbides present in the AM 17-4PH is unclear. Moreover, prior reports have documented fundamental differences in the general deformation behavior between conventional and AM materials135  as well as in hydrogen-charged vs. hydrogen-free precipitation-hardened alloys.117  Whether or not such differences could manifest in the near-crack region for the AM and wrought 17-4 PH is also not known. It should be recognized that an evaluation of the local deformation will not be able to explicitly decouple all of these potential contributions given that they will be acting simultaneously and likely synergistically. However, local near-crack evaluations will provide insights on the extent to which these factors are potential contributors, which could then be leveraged to justify more targeted studies. Such studies of the near-crack deformation for the current materials are presently underway and will be reported in an upcoming manuscript.

Implications

The present study demonstrates a systematic increase in SCC susceptibility for AM 17-4PH relative to wrought 17-4PH with similar yield strength. Leveraging established mechanistic frameworks, the increased susceptibility of the AM material was speculated to arise from five potential factors: (1) differences in grain boundary impurity content, (2) widespread porosity, (3) the presence and/or modification of secondary phase particles, (4) differences in hydrogen-metal interactions between the AM and wrought material, and (5) differences in crack-tip deformation behavior for AM and wrought materials. The latter two contributions will be addressed in future studies, but additional commentary can be provided on the other three possibilities. First, regarding the oxide inclusions, the elemental analysis reported in Table 2 indicates an order of magnitude increase in the O content of the AM material relative to the wrought condition, driven by the high O content of the powder. While the exact contribution of the observed oxide inclusions remains unclear, reducing the O content would minimize such inclusions. Second, grain boundary impurity segregation and Cr23C6 precipitation are likely related to the slow cooling profile associated with the HIP treatment. Given that this cooling rate was an unspecified standard practice for the HIP supplier, the above results and discussion indicate that it would be worthwhile to evaluate the effect of post-HIP cooling rate on SCC behavior in AM 17-4PH steel. In addition to enabling the development of an SCC-resistant HIP protocol, the results of such a study could be used to justify the addition of a specific HIP protocol for 17-4PH in ASTM F3301.45  Lastly, it is clear from microscopy of the metallurgical cross sections (Figure 3) and fracture surfaces (Figure 10) that the sub-μm porosity was resistant to the HIP treatment. As such, it is expected that such small-scale porosity will be present in an AM final component. Given the challenges associated with detecting sub-μm pores using conventional techniques, additional efforts are needed to rigorously understand the influence of these small pores on the SCC susceptibility of 17-4PH. If the proposed deleterious role of small pores on SCC resistance is confirmed, then changes in HIP protocol should be studied with a focus on consistently eliminating or minimizing this sub-μm sized porosity.

Fracture mechanics-based experiments were performed to reveal the relative susceptibilities of wrought and AM forms of 17-4PH to SCC in 0.6 M NaCl as a function of applied electrochemical potential. Intrinsic features leading to differences between the two material forms were evaluated using a combination of compositional, fractographic, and microstructural analysis. The following are key conclusions from this comparative study.

  • AM 17-4PH has similar SCC performance to wrought at open-circuit and anodic potentials in 0.6 M NaCl. Under rising stress intensity loading (), both AM and wrought underwent crack tip blunting with no apparent SCC crack extension with Eapp ∼ EOCP. At anodic potentials (Eapp > EOCP), similar blunting was observed until excessively anodic potentials led to general dissolution in both AM and wrought.

  • At cathodic-applied potentials (Eapp < EOCP), AM 17-4PH consistently underperformed the wrought 17-4PH SCC performance. A combination of rising () and constant () stress intensity testing showed that AM 17-4PH cracked up to 5-fold faster with Eapp susceptibility ranges 200 mV to 400 mV wider than wrought depending on heat treatment.

  • Increased SCC susceptibility in AM 17-4PH was linked to fracture dominated by intergranular cracking morphologies. High-magnification examination of the AM intergranular facets showed evidence of toughness-reducing secondary particles and pores not present in wrought 17-4PH. Based on analysis of 17-4PH phase diagrams, bulk composition analysis results, and the heat treatment protocols used in this study, the particles are likely a combination of Cr23C6 carbides and oxide inclusions.

  • Potential factors responsible for the increased susceptibility of the AM 17-4PH were informed by a decohesion-based model for hydrogen-assisted cracking and include: (1) differences in grain boundary impurity content, (2) widespread porosity, (3) the presence and/or modification of secondary phase particles, (4) differences in hydrogen-metal interactions between the AM and wrought material, and (5) differences in crack-tip deformation behavior for AM and wrought materials.

  • While additional studies are needed to quantify hydrogen-metal interactions and near-crack deformation behavior in the current materials, it is likely that the (1) significant O content (600 wppm) in the powder, (2) slow cooling (310°C/h) during the HIP process, and (3) argon gas used as a powder atomizing media and as a build atmosphere resulting in widespread sub-μm porosity all negatively affected the SCC resistance of AM 17-4PH. Such initial insights can be leveraged to inform targeted processing modifications to potentially improve the SCC behavior of AM 17-4PH.

(1)

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.

Trade name.

This research was financially supported by the Office of Naval Research under Grant No. N00014-18-1-2427 with Dr. Airan Perez as the Scientific Officer. Helpful discussions with Prof. John Scully and Ms. Lauren Singer at the University of Virginia are gratefully acknowledged.

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