Effect of Nitrogen Content on the Cavitation Erosion Resistance of 316LN Stainless Steel

Abstract

Cavitation erosion is a predominant failure mode of austenitic stainless steels in corrosive fluid environments, severely limiting their durability in nuclear piping and hydraulic components. In this study, five 316LN steels with 0.008–0.34 wt.% nitrogen content were fabricated, and both short-term (2 h) and long-term (24 h) cavitation tests were performed to elucidate the effect and mechanism of nitrogen. Increasing nitrogen markedly enhanced cavitation resistance: after 24 h, the cumulative mass loss decreased by 36%, 52%, 60%, and 71% for 09N, 17N, 22N, and 34N relative to 00N, accompanied by lower surface roughness, shallower pit depth, and a prolonged incubation stage. SEM revealed a progressive damage process from twin/high-angle grain boundaries to intragranular deformation bands and finally to spalling at slip intersections, whereas high-N steels exhibited only slight local detachment. TEM demonstrated that nitrogen transformed dislocations from random networks into dense slip bands and planar arrays with stacking faults, raising hardness from ~140 HV to ~260 HV. EBSD further confirmed strain-induced martensite transformation under severe deformation, providing additional strengthening. These results reveal that nitrogen improves cavitation resistance by tailoring dislocation structures and enhancing strength–plasticity compatibility, offering guidance for the design of high-performance austenitic stainless steels in cavitation environments.

1. Introduction

316LN is a Cr–Ni–N austenitic stainless steel developed by introducing nitrogen into 316L. Cr promotes the formation of a dense Cr-rich passive film and thus significantly enhances corrosion resistance [1]. Ni stabilizes the austenitic structure and improves toughness and ductility [2], while nitrogen, as an interstitial element, not only expands and stabilizes the austenite phase field but also strengthens the steel through solid-solution strengthening and improves work-hardening capability [3]. Owing to its excellent corrosion resistance and mechanical properties, 316LN has been widely used in nuclear power pipelines [4], petrochemical equipment [5], and marine valves [6], where both corrosion resistance and load-bearing capability are required.

However, in high-speed flow environments containing corrosive media, 316LN stainless steel remains susceptible to damage due to cavitation phenomenon. Cavitation typically occurs in regions with flow disturbances, such as local contractions and welds, where the repeated formation and collapse of bubbles generate high-frequency micro-jets and pressure waves that induce severe plastic deformation, erosion wear, fatigue crack initiation and propagation, and ultimately pit formation [7,8,9]. During long-term service, components such as nuclear pipelines, turbine blades, and pump housings face similar cavitation risks. Previous studies have shown that austenitic stainless steels exhibit better cavitation resistance than some copper alloys and high-strength steels due to their high plasticity and energy absorption capacity [10]. In addition, alloying elements play a crucial role in cavitation resistance; for example, Mo enhances resistance to pitting corrosion and delays pit propagation under cavitation [11]. Nitrogen has been reported to improve both strength and corrosion resistance in austenitic stainless steels [12] and has shown potential to enhance cavitation resistance [13]. However, excessive nitrogen may lead to cellular Cr₂N precipitation [14], resulting in local Cr depletion and degradation of corrosion and mechanical properties. Although the effects of nitrogen on creep [15], corrosion behavior [16], and hot deformation behavior [17] have been relatively well studied, a systematic understanding of its influence on cavitation resistance and the underlying mechanisms in 316LN stainless steel remains limited. Furthermore, whether cavitation involves electrochemical corrosion remains controversial [18,19]. It is noted that in chloride-containing environments, cavitation damage may couple with electrochemical dissolution, leading to a synergistic acceleration of material removal. Recent studies have shown that bubble collapse can disrupt passive films and enhance localized corrosion, while corrosion can in turn facilitate pit initiation and crack propagation under repeated micro-jet impacts. Although higher nitrogen levels (>0.4 wt.%) are technically achievable, such compositions readily promote Cr₂N precipitation during solidification or subsequent thermal exposure, resulting in local chromium depletion and degradation of corrosion resistance and mechanical stability [20].

Therefore, five 316LN steels with nitrogen contents ranging from 0.008 to 0.34 wt.% were subjected to short-term and long-term cavitation tests, combined with surface morphology and microstructural characterization, to systematically elucidate the influence of nitrogen on cavitation resistance and the associated mechanisms. The results provide guidance for the design and application of cavitation-resistant austenitic stainless steels in nuclear, hydraulic, and marine environments.

2. Materials and Methods

2.1. Materials

Five 316LN austenitic stainless steels were produced by means of melting in a pressurized induction furnace followed by pressurized electroslag remelting to ensure high purity and nitrogen solubility. These chemical compositions of the five 316LN austenitic stainless steels were measured. Nitrogen content was determined using an ON/H/N elemental analyzer (LECO ON836, LECO Corporation, St. Joseph, MI, USA) with three repeated measurements to ensure accuracy. The concentrations of metallic elements (Fe, Cr, Ni, Mo, Mn, etc.) were analyzed via inductively coupled plasma optical emission spectroscopy (ICP-OES, Agilent Technologies, Santa Clara, CA, USA). The measured compositions confirmed that the nitrogen contents of the five steels were 0.008, 0.09, 0.17, 0.22, and 0.34 wt.%, respectively (denoted 00N, 09N, 17N, 22N and 34N). The chemical compositions after electroslag remelting are shown in

Table 1. Chemical composition of 316L(N) (wt.%).

	N	C	Cr	Ni	Mo	Si	S	P	Fe
00N	0.008	0.015	18.2	12.47	2.52	0.47	0.0055	0.013	Bal.
09N	0.09	0.02	18.15	12.36	2.51	0.49	0.0044	0.015	Bal.
17N	0.17	0.018	18.3	12.55	2.52	0.45	0.0039	0.011	Bal.
22N	0.22	0.018	18.23	12.45	2.52	0.39	0.0045	0.019	Bal.
34N	0.34	0.02	18.27	12.57	2.51	0.75	0.005	0.011	Bal.

. The ingots were cut into slabs (30 mm × 30 mm × 30 mm) and hot-rolled at 1100 °C to a total reduction of 60%. The reduction was achieved in four passes, with the thickness decreased by ~20%, ~15%, ~15%, and ~10% in successive passes. All specimens were subjected to solution treatment to ensure a fully austenitic microstructure and comparable initial grain size. The detailed solution-treatment temperatures and holding times for each nitrogen level are listed in

Table 2. Solution-treatment regimes and average grain sizes of 316L(N).

Alloys	00N	09N	17N	22N	34N
Regimes	1150 °C/ 150 min	1200 °C/ 30 min	1200 °C/ 70 min	1200 °C/ 80 min	1150 °C/ 95 min
D(μm)	112 ± 2	110 ± 3	112 ± 3	115 ± 1	107 ± 5

.

2.2. Cavitation Erosion Tests

Cavitation tests were performed on solution-treated specimens using an XOQS-2500 (XOQS-2500, Xianou Instrument Equipment Co., Ltd., Nanjing, China) ultrasonic cavitation apparatus. Disk specimens (⌀10 mm × 3 mm) were prepared by EDM (Electrical Discharge Machining, EDM, commercial equipment, Suzhou, China), ground to 2000 grit and polished with 2.5 μm diamond suspension. The horn tip diameter was 15.9 mm; the vibration frequency and amplitude were 20 kHz and 20 μm, respectively. Tests were conducted in intermittent mode (10 s cycle: 9.5 s cavitation, 0.5 s pause). Effective acoustic power was verified by calorimetry and reported as power density. The horn–specimen gap r (0.5 mm) was calibrated before each series. Deionized water was used as the test medium, and the water was replaced every 8 h. The temperature was controlled at 25 ± 1 °C using a recirculating chiller with feedback and logged throughout the tests. Dissolved oxygen was monitored with a portable dissolved oxygen (DO) meter and maintained near the air-saturation value at 25 °C. Specimens were centered under the horn with a positioning jig to ensure coaxial alignment; the in-plane offset was limited to ≤ 0.1 mm. The disk geometry was reproducible within ±0.1 mm in diameter and thickness across batches, ensuring test comparability. These cavitation test parameters were selected according to ASTM G32-2016 [21] and are consistent with commonly used conditions in cavitation studies on austenitic stainless steels, ensuring comparability and reproducibility of the results [22].

Two sets of tests were performed. The first was a 24 h mass-loss test. Three specimens were tested for each composition. During the initial incubation stage, mass was measured every 1 h to determine the onset of material loss. After mass loss occurred, measurements were taken every 3 h and the average cumulative mass loss was plotted. Mass loss was calculated as Equation (1):

Δm = (m₀ − m₁)/S

(1)

where m₀ is the initial mass (mg), m₁ is the mass after cavitation (mg), and S is the exposed area (cm²). The second set consisted of 2 h short-term cavitation tests.

2.3. Microstructural and Surface Characterization

Initial grain sizes were examined using an Axio Observer A1m optical microscope (Carl Zeiss, Jena, Germany). Specimens were etched in 50% aqua regia (30 mL HCl + 10 mL HNO₃ + 40 mL H₂O) for 3 min prior to metallography. SEM observations were performed on a Hitachi S-3400 (Hitachi High-Technologies, Tokyo, Japan) at 3 kV using an SE2 detector. Three-dimensional surface profiles, roughness (Ra) and pit depths were measured using a VEECO Contour GT-K1 optical profiler (Veeco Instruments Inc., Plainview, NY, USA. scan area 1.25 × 1.00 mm) and analyzed with Vision software (Version 6.5, Bruker Nano Surfaces, Tucson, AZ, USA).

To analyze microstructural evolution during deformation, transmission electron microscopy (TEM, FEI Titan ETEM G2, FEI Company, Hillsboro, OR, USA) was performed on specimens. Prior to TEM observation, the specimens were subjected to ex-situ uniaxial tensile deformation to a true strain of 2.5% at room temperature. This strain level was chosen to introduce initial dislocation structures while minimizing deformation twinning and suppressing strain-induced martensite formation, allowing the intrinsic effects of nitrogen content on deformation behavior to be clearly revealed. Disks (0.3 mm thick) were cut by electrical discharge machining, ground to 40 μm (2000# and 4000# papers), and twin-jet electropolished in acetic acid + 5% perchloric acid at 20 V (flow rate 25) until a transparency threshold of 80 was reached. Imaging was performed in bright-field mode under two-beam conditions to enhance dislocation and stacking-fault contrast; selected g-vectors (g = (111)) were used. Multiple representative fields and magnifications were examined for each alloy to ensure reproducibility.

Electron backscatter diffraction (EBSD) was used to characterize the phase map and strain-induced martensite in specimens with 50% deformation. EBSD analyses were conducted on tensile-deformed specimens (50% strain) using a step size of 0.1 µm and an accelerating voltage of 20 kV. Noise reduction was applied through grain dilation and neighbor orientation averaging, and a minimum area of five pixels was set for phase assignment. These parameters ensured reliable indexing and minimized artifacts. Specimens were electropolished in 90% ethanol + 10% perchloric acid at 26 V for 1 min. EBSD mapping was performed with a step size of 0.1 μm, and data were analyzed using TSL OIM Analysis™ software (EDAX, A Division of AMETEK, Mahwah, NJ, USA).

To establish the relationship between cavitation resistance and hardness, Vickers microhardness of the solution-treated specimens was measured using an FM-ARS-9000 automatic system. A load of 25 gf was used for Vickers microhardness testing to ensure consistent measurement conditions among the specimens. Microhardness was measured using a Vickers tester under a load of 25 gf with a dwell time of 10 s. At least ten indentations were taken per specimen, and the average value was reported.

3. Results

3.1. Cavitation Mass Loss

Figure 1a presents the mass-loss curves of the 316LN steels with different nitrogen contents after 24 h of cavitation. With increasing exposure time, all specimens experienced an incubation period, a rapid mass-loss stage, and finally a steady state. Meanwhile, mass loss gradually decreased with increasing nitrogen content. The cumulative mass loss reached 11.3 mg·cm⁻² for 00N, 7.2 mg·cm⁻² for 09N, 5.4 mg·cm⁻² for 17N, 4.5 mg·cm⁻² for 22N, and only 3.3 mg·cm⁻² for 34N, corresponding to reductions of 36%, 52%, 60%, and 71%, respectively, compared with 00N.

Figure 1b shows the relationship between cumulative mass loss and nitrogen content. An exponential decrease in mass loss with increasing nitrogen content was observed, indicating that nitrogen effectively enhances the cavitation resistance of 316LN steel. The relationship between cumulative mass loss and nitrogen content was fitted using a nonlinear least-squares exponential regression. The fitted equation is:

$${\Delta m}_{\left(\mathrm{N}\right)}=2.6+9.2\times {6.7}^{\mathrm{N}}(\mathrm{mg}·{\mathrm{cm}}^{-2})$$

(2)

with R² = 0.982 and a mean square error (MSE) of 0.19 (mg·cm⁻²)², indicating a high fitting accuracy. This regression is applicable within the nitrogen content range of 0.008–0.34 wt.% used in this study. It should be emphasized that the exponential trend describes the experimental behavior in this composition range and does not suggest the existence of a theoretical minimum mass loss.

3.2. Three-Dimensional Surface Morphology and Roughness After Cavitation

Figure 2 shows the 3D surface morphology and the corresponding histograms of depth distribution of the 316LN steels after 24 h of cavitation. Each 3D image covers an area of 1.25 mm × 1.0 mm, corresponding to the region of most severe cavitation damage. The color scale indicates pit depth (−80 μm to 0 µm).

As the nitrogen content increased, both the number and depth of cavitation pits decreased markedly. In Figure 2(a1), the 00N specimen exhibited numerous, uniformly distributed pits and a high surface roughness. The depth distribution in Figure 2(a2) revealed the average depth of ~61.7 µm. For 09N (Figure 2(b1)), the number of pits was significantly reduced and their distribution became non-uniform, resulting in a lower surface roughness. The depth distribution (Figure 2(b2)) showed an average depth of ~45.0 µm. With further nitrogen addition, the 17N specimen exhibited fewer deep pits and a smoother surface (Figure 2(c1)), and the average depth decreased to ~13.1 µm (Figure 2(c2)). The surface morphology and average depth of the 22N specimen remained similar to that of 17N (Figure 2(d1,d2)). However, the 34N specimen exhibited extremely fine and uniformly distributed pits, with the average distribution of only ~4.6 µm (Figure 2(e1,e2)).

Figure 3 illustrates the relationship between surface roughness and nitrogen content after cavitation. Surface roughness (Ra) was calculated from the entire scanned area using Vision software. It can be observed that the Ra decreases markedly with increasing nitrogen content. Curve fitting revealed the following empirical relationship between Ra and N, which is given as Equation (3).

$$\mathrm{Ra}=1.6+3.9 \times {2.2}^{\mathrm{N}}\left(\mathsf{\mu }\mathrm{m}\right)$$

(3)

where Ra denotes surface roughness and N represents nitrogen content. The regression was obtained using a least-squares fitting method. The fitting quality was evaluated by the coefficient of determination (R² = 0.99) and mean square error (MSE = 0.27). This empirical equation is applicable within the nitrogen range of 0.008–0.34 wt.%. A minimum Ra value was not reached in this range, as the surface roughness continued to decrease with increasing nitrogen content. Notably, both cumulative mass loss and surface roughness exhibit an exponential dependence on nitrogen content, indicating that increasing nitrogen plays a significant role in enhancing the cavitation resistance of 316LN steel.

3.3. Short-Term Cavitation Damage Behavior and Mechanism

Figure 4 presents the SEM surface morphologies of 316LN steels with different nitrogen contents after 2 h of cavitation.

Overall, even short-term cavitation can induce noticeable plastic/fatigue damage on the specimen surfaces, and the damage characteristics vary significantly with nitrogen content. For the 00N specimen (Figure 4(a1)), numerous cavitation pits and spalling were already observed, with pits gradually connecting to form a highly rough surface. The magnified image (Figure 4(a2)) revealed sharp pit edges and severe damage. For 09N, extensive strip-like spalling appeared after just 2 h, extending along twin boundaries and slip bands, with severe damage at multiple slip band intersections (Figure 4(b1)). The magnified view showed strips with widths of several tens of micrometers; at the intersections, clear fracture steps, warped edges, and fine cracks were visible, indicative of shear-dominated failure morphology (Figure 4(b2)). This suggests that under short-term micro-jet impacts, low-nitrogen steel is more prone to dislocation accumulation and spalling. For 17N, the strip-like spalling was significantly reduced. The low-magnification image showed a decrease in overall roughness and interrupted connectivity of large spalling regions (Figure 4(c1)). The magnified view revealed small pits of ~4 µm and relatively sharp pit edges, with only a few deep spalling sites at intersecting slip bands; slight warping along the slip bands was also visible (Figure 4(c2)).

For 22N, twin and grain boundary contours were distinct, and the overall surface integrity was further enhanced. Only a few dispersed shallow depressions and ripple-like textures were observed in the low-magnification image (Figure 4(d1)). The magnified view showed discontinuous slip bands and minor micro-cracks, with almost no observable spalling steps or large pit cores, indicating that this composition can better maintain surface continuity under short-term high-frequency impacts, and local strain manifests primarily as microplastic fluctuations (Figure 4(d2)). For 34N, after 2 h of cavitation, most areas remained nearly undamaged, with only slight spalling or shallow pits at a few twin boundaries. The magnified image revealed minor damage at occasional intersecting deformation bands within twins, but no fatigue striations or obvious fracture steps were observed.

Observations and analyses of short-term cavitation microstructures indicate that the cavitation resistance of 316LN steel improves significantly with increasing nitrogen content. The underlying damage mechanism can be described as follows: micro-jet impacts from collapsing bubbles initially induce fluctuations and protrusions along twin boundaries, triggering spalling. With continued cavitation, damage gradually extends into twins, with intersecting deformation bands forming minor spalling at the intersections. As cavitation progresses, grain boundaries begin to protrude and spall, activating limited slip systems; these slip systems undergo micro-jet impact, warp, and develop edge spalling, followed by activation of multiple slip systems. Severe spalling occurs at intersecting slip bands, and the spalled regions gradually expand until the surface layer is completely removed.

3.4. Surface Morphology After 24 h Cavitation

Figure 5 shows the SEM surface morphology of the 316LN steels after 24 h of cavitation. The 00N specimen exhibited severe damage, with large spalling regions and deep interconnected pits forming a highly rough and non-uniform surface (Figure 5(a1)). In the magnified view (Figure 5(a2)), distinct fracture steps, fatigue striations and sharp pit edges were observed, indicating cumulative impacts and mixed brittle–ductile fracture. Residual surface layers prior to spalling were still visible, exposing underlying surfaces and suggesting multiple cycles of material removal. This severe morphology is consistent with its highest cumulative mass loss, confirming that 00N has the poorest cavitation resistance.

Compared with 00N, the 09N surface showed noticeably reduced roughness (Figure 5(b1)), with fewer large spalling pits and only medium-depth irregular depressions. The magnified view (Figure 5(b2)) still revealed sharp pit edges and ~4 µm micro-pits, indicating that early spalling was still active but damage became more dispersed. The 17N specimen exhibited a smoother surface (Figure 5(c1)), with significantly fewer and shallower pits than 09N. Large spalling almost disappeared, and only slight pit nuclei were observed locally. In the magnified image (Figure 5(c2)), pit edges appeared rounded and fracture features shifted from cleavage to plastic deformation marks, with no obvious steps or fatigue striations, indicating delayed crack propagation and improved cavitation resistance. The 22N surface remained largely intact (Figure 5(d1)), with almost no deep pits or large spalling. Damage was mainly expressed as fine ripple- or layer-like deformation. Compared to 17N, pit features were further weakened and more uniformly distributed at smaller scales. The magnified view (Figure 5(d2)) showed continuous undulating structures without fracture steps, with only minor local spalling, suggesting stronger resistance to impact fatigue and efficient energy dissipation under cyclic micro-jets.

The 34N specimen maintained the most intact surface (Figure 5(e1)), exhibiting a smooth and dense layer with almost no pits or spalling, and even some areas remained unpeeled. Only very shallow micro-depressions were observed locally. The high-magnification image (Figure 5(e2)) showed only slight ripple-like or plastic undulations, without fatigue striations, fracture steps, or crack sources, and damage was minimal, localized, and slow to develop. Compared with 00N, these results clearly demonstrate that increasing nitrogen effectively suppresses pit initiation and delays cumulative mass loss, allowing the surface to remain stable even under prolonged cavitation exposure.

4. Discussion

Cavitation damage is generally caused by repeated micro-jet impacts and fatigue accumulation. The process essentially involves local plastic deformation, crack initiation, and surface spalling under high-frequency loading. Therefore, the material’s strength level, plasticity coordination, and phase interface stability are considered key factors determining cavitation resistance [23].

In this study, with increasing nitrogen content, both the surface roughness and cumulative mass loss of 316LN steel decreased significantly (Figure 1, Figure 2 and Figure 3), indicating that high-nitrogen specimens are more resistant to cavitation-induced pit formation and spalling. This suggests that nitrogen may enhance cavitation resistance by altering the mechanical response of the material.

Figure 6 shows the relationship between microhardness and nitrogen content for solution-treated 316LN steels with five different nitrogen levels. It can be seen that the hardness of 316LN steel increases significantly with nitrogen content, which contributes to the improvement of its cavitation resistance. The relationship between nitrogen content and hardness was fitted using a nonlinear least-squares regression approach. The fitted equation is expressed as:

$$\mathrm{HV}=349{\mathrm{N}}^2+80\mathrm{N}+137 (\mathrm{N}/{\mathrm{mm}}^2)$$

(4)

with a high goodness-of-fit (R² = 0.99) and a mean square error of approximately 12 HV². This relationship is applicable within the nitrogen range of 0.008–0.34 wt.%, over which hardness increases continuously with nitrogen content.

It should be noted that the hardness values shown in Figure 6 represent the initial microhardness of the solution-treated steels, measured prior to cavitation. No significant hardness gradient was expected across the surface layer under the current test conditions, as cavitation-induced work hardening is typically confined to a depth below a few micrometers. Therefore, the improvement in cavitation resistance is mainly attributed to the intrinsic strengthening effect of nitrogen through solid-solution hardening and dislocation structure modification.

To further investigate the mechanism by which nitrogen content affects the mechanical properties of 316LN stainless steel, five specimens with different nitrogen levels were subjected to low-strain tensile deformation. Figure 7 shows representative TEM micrographs of 316LN steels under 2.5% strain.

The results indicate that in the low-nitrogen specimen (00N), dislocations were distributed in a relatively disordered manner, and dislocation interactions were easily annihilated (Figure 7a). With increasing nitrogen content (09N→17N→22N), More slip bands and dislocation accumulations appeared both within grains and along grain boundaries, which acted as the primary sites for dislocation pile-up and stress concentration. Planar slip arrays also began to emerge (Figure 7b–d). In the high-nitrogen specimen (34N), numerous planar slip arrays (Figure 7e) and stacking faults (SFs) (Figure 7f) were observed across different fields.

Previous studies have shown that a reduction in stacking fault energy can induce short-range ordering within the austenitic matrix, promoting planar slip and stacking fault formation [24,25,26], thereby significantly increasing the slip activation stress and dislocation pinning effect [27,28], which in turn enhances the yield strength. Therefore, it can be concluded that nitrogen substantially increases the yield strength and hardness of the material by regulating dislocation motion and slip behavior.

Although nitrogen could not be directly detected using EDS due to its low atomic number, the TEM observations in Figure 7 confirm the absence of Cr₂N precipitation, indicating that nitrogen remains dissolved interstitially in the austenitic matrix. Previous studies have shown that interstitial nitrogen exhibits strong interaction with Cr atoms, promoting the formation of short-range order (SRO) clusters such as Cr–N and Cr–Mo–N [29,30,31]. The presence of SRO reduces the stacking fault energy (SFE), which increases the likelihood of planar slip and stacking-fault formation [25,32]. Correspondingly, the TEM results in high-nitrogen steels (Figure 7e,f) show well-developed planar slip bands and extended stacking faults, as opposed to the more random dislocation configurations observed in low-nitrogen steels. In addition, nitrogen interacts electrostatically with dislocation cores and forms Cottrell atmospheres, which increases the stress required to initiate dislocation motion [33,34]. These effects collectively raise the resistance to localized plastic collapse under micro-jet impact, thereby improving cavitation resistance.

In addition, austenitic stainless steels with low stacking fault energy can undergo strain-induced martensitic (α′-M) transformation under large plastic deformation. During cavitation, the micro-jet strain rates generated by ultrasonic bubble collapse can reach 10³–5 × 10⁶ s⁻¹, with pressures of several thousand bars [35,36]. However, due to the rough surface of cavitation specimens, EBSD observation is not directly feasible. Therefore, this study examined the phase map in specimens subjected to large plastic deformation, as shown in Figure 8.

In these maps, austenite and α′-M are distinguished by color contrast, and the martensite fraction was estimated based on the area percentage of the martensitic phase. The results indicate that all specimens underwent strain-induced martensitic transformation, and the martensite fraction increased with nitrogen content, rising from approximately 8.6% in 00N steel (Figure 8a) to 10.2% in 17N steel (Figure 8b) and 15.8% in 34N steel (Figure 8c). Since martensite has a higher intrinsic strength, the increased martensite fraction contributes to localized strengthening and improves resistance to micro-jet impact and surface layer collapse [37,38,39].

However, it should be noted that in low-nitrogen specimens, the hardness difference between γ and α′-M phases is large, leading to severe stress concentration at the interface. The α′-M phase is therefore more prone to brittle spalling, producing flat fracture surfaces. In high-nitrogen specimens, the matrix is harder, reducing the strength difference between the matrix and α′-M phase, resulting in more coordinated phase transformation. Crack propagation is suppressed, and spalling is delayed. This indicates that nitrogen not only increases overall strength but also optimizes the strength gradient and interface stability between the matrix and martensite, thereby improving resistance to fatigue spalling. This also explains why the incubation period of cavitation gradually increases with nitrogen content (Figure 1).

It is noteworthy that in this study, ultrapure water was used as the medium, and no obvious corrosion products were observed in the numerous SEM images, indicating that the damage was primarily driven by mechanical impact and fatigue. This does not contradict other reports on cavitation–corrosion coupling effects [11], as the cavitation response is highly dependent on both the material and the medium. In this study, 316LN exhibited good passivation, and the chloride ion concentration was extremely low; therefore, mechanical damage was dominant. Nitrogen, by strengthening the matrix and regulating the microstructure, demonstrated an enhanced cavitation resistance under purely mechanically driven conditions.

In summary, the improvement of cavitation resistance by nitrogen content is not a result of a single factor but arises from multiple synergistic mechanisms, including increased strength, modification of dislocation structures, promotion of planar slip pile-ups, coordination of strain-induced martensitic transformation, and optimization of interface stability. This mechanistic chain, progressing from macroscopic properties through microscopic dislocations to phase transformation behavior, provides a systematic theoretical basis for compositional optimization and cavitation-resistant design of 316LN steel.

It should be noted that all experiments were performed in deionized water under specific ultrasonic cavitation conditions. Therefore, the results primarily reflect the mechanical–fatigue-driven cavitation behavior of 316LN steels in the absence of corrosion coupling.

5. Conclusions

Cavitation tests were performed on 316LN stainless steels with five nitrogen contents to study the effects of nitrogen on cavitation resistance. The main findings are as follows:

(1)

After 24 h of cavitation, cumulative mass loss decreased significantly with nitrogen content. Compared to 00N, 09N, 17N, 22N, and 34N steels showed reductions of 36%, 52%, 60%, and 71%, respectively. Surface roughness also decreased exponentially with increasing nitrogen.

(2)

Cavitation damage initiated at twin boundaries and high-angle grain boundaries and then propagated into twin interiors and intragranular deformation bands.

(3)

Microhardness increased from ~140 HV (00N) to ~260 HV (34N). TEM showed that dislocations evolved from disordered to dense slip bands and planar slip arrays, with stacking faults forming, indicating that nitrogen strengthens the material and enhances resistance to impact deformation.

(4)

The results indicate that nitrogen exists as interstitial atoms forming short-range ordered regions, which reduce SFE and promote planar slip, thereby contributing to the enhanced cavitation resistance.