Interface velocity-driven non-equilibrium nitrogen supersaturation in additive manufacture: a universal strategy for breaking strength-corrosion trade off

Interface velocity-driven non-equilibrium nitrogen supersaturation in additive manufacture: a universal strategy for breaking strength-corrosion trade off | Research Square window.SnipcartSettings = { analytics: { enabled: false } }; (function() { var accessVector = localStorage.getItem('access_vector') || ''; window.dataLayer = window.dataLayer || []; if (accessVector) { window.dataLayer.push({ user: { profile: { profileInfo: { snid: accessVector } } } }); } })(); (function(w,d,s,l,i){w[l]=w[l]||[];w[l].push({'gtm.start':new Date().getTime(),event:'gtm.js'});var f=d.getElementsByTagName(s)[0],j=d.createElement(s),dl=l!='dataLayer'?'&l='+l:'';j.async=true;j.src='https://www.googletagmanager.com/gtm.js?id='+i+dl;f.parentNode.insertBefore(j,f);})(window,document,'script','dataLayer','GTM-K279D39R'); Browse Preprints In Review Journals COVID-19 Preprints AJE Video Bytes Research Tools Research Promotion AJE Professional Editing AJE Rubriq About Preprint Platform In Review Editorial Policies Our Team Advisory Board Help Center Sign In Submit a Preprint Cite Share Download PDF Article Interface velocity-driven non-equilibrium nitrogen supersaturation in additive manufacture: a universal strategy for breaking strength-corrosion trade off Mujin Yang, Daobin Zhang, Dingding Zhu, Bo Du, Minglin He, Xiaomei Zhang, and 7 more This is a preprint; it has not been peer reviewed by a journal. https://doi.org/ 10.21203/rs.3.rs-9249060/v1 This work is licensed under a CC BY 4.0 License Status: Posted Version 1 posted You are reading this latest preprint version Abstract Achieving synergistic high strength and corrosion resistance in high-nitrogen stainless steels remains challenging due to inherent trade-offs: dispersion strengthening enhances strength, but heterogeneous interfaces often act as micro-galvanic corrosion or pitting nucleation sites, compromising corrosion resistance. Laser powder bed fusion (LPBF) with its extreme thermal cycles, offers a non-equilibrium strategy to address this limitation. Here, using multi-scale microstructural characterization and thermo-kinetic simulations methods, we report that during melt pool solidification, interface velocity-regulated kinetics trap excess nitrogen in the ferritic matrix, inducing non-equilibrium nitrogen supersaturation. Subsequently during fast cooling process, this supersaturated nitrogen combines with chromium to form dense superfine CrN lamella. Coupled with fine grain structures from rapid solidification, this two-stage non-equilibrium process achieves a synergy of high tensile strength (~1700 MPa) and pitting resistance (~1000 mV), breaking the classical strength-corrosion trade-off. Key mechanisms include: (1) superfine CrN lamellae minimizing adjacent Cr-depleted zones, (2) grain refinement suppressing precipitate-free zones (PFZ), and (3) nitrogen-stabilized grain boundaries enhancing corrosion resistance. By controlling interface velocity in LPBF non-equilibrium solidification, we establish a nitrogen supersaturation pathway that tailors hierarchical microstructures (grains, boundaries, nanoprecipitates), resolving the strength-corrosion trade-off in additively manufactured high-nitrogen alloys. Physical sciences/Materials science/Structural materials/Metals and alloys Physical sciences/Materials science/Structural materials/Mechanical properties Figures Figure 1 Figure 2 Figure 3 Figure 4 Figure 5 Introduction Laser powder bed fusion (LPBF) has unlocked new frontiers in metallurgy by leveraging its characteristic extreme thermal cycles—featuring rapid solidification (~ 10 3 –10 6 K/s) and repeated remelting—to create non-equilibrium microstructures and defect landscapes unattainable through conventional processing [ 1 , 2 ]. This ability to dictate non-equilibrium phase transformations presents a compelling opportunity to revisit and potentially solve persistent issues in alloy design, including the control of detrimental secondary phase precipitation in systems like high nitrogen duplex stainless steels (DSSs) [ 3 , 4 ]. High nitrogen DSSs with their balanced ferrite-austenite microstructure, are prized for their excellent combination of toughness and corrosion resistance, serving critical roles in marine, chemical, and energy industries [ 5 , 6 ]. However, their performance is intrinsically constrained by a classic trade-off: conventional strengthening methods, such as cold working or precipitation hardening via secondary phases, often degrade their corrosion resistance. A central challenge lies in controlling the precipitation of chromium nitrides (Cr 2 N/CrN) [ 7 , 8 ]. During conventional processing, these nitrides typically precipitate in the temperature range of 700–1000°C, either during slow cooling or isothermal aging. Even under rapid quenching, so-called “quenched-in” nitrides can form within the ferrite matrix, as the high cooling rate suppresses nitrogen diffusion into austenite, leading to supersaturation and subsequent intragranular precipitation [ 9 , 10 ]. While refining the nitride morphology, these precipitates still deplete chromium in adjacent areas and are widely reported to impair pitting corrosion resistance. Therefore, avoiding or minimizing nitride formation has been a guiding principle in the thermo-mechanical processing of high nitrogen DSSs. In pursuit of this goal, several strategies have been explored to mitigate this trade-off. Conventional approaches often focus on precise thermo-mechanical processing routes to control the nucleation and growth of deleterious phases [ 11 , 12 ]. Alloy design innovations, such as optimizing the Cr/Ni/Mn/N ratio or adding nitride-stabilizing elements like Cu, have also been attempted to refine precipitate distribution [ 13 – 15 ]. More recently, advanced manufacturing techniques like directed energy deposition (DED) have been employed to create graded microstructures, yet control over nanoscale precipitation remains challenging [ 16 ]. While these methods have met with varying degrees of success, they often involve complex post-processing or still struggle to completely decouple the strength enhancement provided by precipitates from their detrimental effect on corrosion resistance. The fundamental challenge lies in the inherent equilibrium or near-equilibrium nature of these processes, which favors the formation of coarse, chromium-depleting precipitates. The LPBF process, with its inherent non-equilibrium conditions, challenges this conventional wisdom. The extreme undercooling and thermal stress inherent to LPBF generate a high density of crystal defects, particularly subgrains, dislocations and vacancies [ 17 – 20 ]. In other alloy systems fabricated by LPBF, such as 316L stainless steel and Inconel 718, these processing-induced defects have been shown to act as potent nucleation sites for secondary phases—a phenomenon termed defect-assisted precipitation [ 21 – 23 ]. This raises a thought-provoking possibility: Can the unique extreme undercooling and special defect distribution induced by thermal stress in LPBF be leveraged to regulate the non-equilibrium supersaturation of nitrogen in the ferrite phase, guide nitride precipitation into dispersed, fine, and favorable morphologies, thereby enhancing strength without compromising corrosion resistance? Specifically, as a key factor determining the solidification rate, the solid-liquid interface migration rate—controlled by laser scanning speed—may serve as an effective means to regulate this supersaturation behavior. However, for high-nitrogen stainless steel processed by LPBF, the influence of laser scanning speed on the microstructure-property relationship remains to be explored. Herein, we demonstrate that LPBF effectively overcomes the traditional strength-corrosion trade-off in a novel high-nitrogen DSSs. We reveal that interface velocity-driven non-equilibrium nitrogen supersaturation, precisely controlled by laser scan speed, promotes the formation of a high dense coherent CrN lamella and substantial grain refinement, concurrently narrowing the precipitation-free zone. Multi-scale characterization shows these microstructural features synergistically lead to a remarkable ~ 400 MPa increase in yield strength, reaching ultra-high strength level, while maintaining exceptional pitting corrosion resistance (E pit ≈ 1000 mV). This work not only provides a new pathway for developing high-performance high-nitrogen stainless steels via additive manufacturing but also offers fundamental insights into phase transformation control under extreme processing conditions. Results Non-equilibrium nitrogen supersaturation strategy in additive manufacturing Figure 1 conceptually illustrates how LPBF regulates non-equilibrium nitrogen partitioning between ferrite and austenite phases, contrasting it with traditional casting. At a constant nitrogen concentration (0.63 wt.%), traditional casting (Fig. 1 a) involves quasi-equilibrium solidification, where nitrogen diffuses to achieve near-equilibrium, resulting in significant nitrogen enrichment in austenite and limited, coarsened nitride precipitation at grain boundary (GB) and other defect sites. In contrast, the AM melt pool (Fig. 1 b) operates under laser scanning speed (SP)-modulated non-equilibrium solidification kinetics, characterized by significantly much higher solidification and cooling rates, even at low scanning speed (LSP) condition (Fig. 1 b 1 ). This rapid process hinders nitrogen diffusion, leading to a non-equilibrium state where ferrite is nitrogen-supersaturated (see point B as arrow pinpointed in nitrogen concentration legend, Fig. 1 ) and austenite is slightly nitrogen-deficient (point B'). During subsequent rapid cooling, CrN nucleates and grows. Under LSP (Fig. 1 b 1 ), relatively large grains (tens of micrometers) form with fine, dispersed CrN lamella, often accompanied by a distinct precipitation-free zone (PFZ) near the austenite/ferrite interface. Increasing scanning speed (HSP in Fig. 1 b 2 ) further elevates solidification rates, inducing significant grain refinement (< 10 µm) and leading to finer, more dispersed CrN precipitates. This combined refinement significantly narrows or eliminates the PFZ, representing a key microstructural difference between LSP and HSP conditions. Regulating non-equilibrium nitrogen supersaturation via laser scan speed Microstructural characterization results in Figs. 2 , 3 , S2 and S3 reveal the impact of SP on nitrogen saturation. IPF maps ( Fig. 2a 1 and 2b 1 ) show no significant texture differences, but increasing SP significantly refines grains. Grain size statistics ( Fig. S2g , S2h ) confirm average grain size reduction from 26 ± 6 µm (LSP) to 6 ± 3 µm (HSP), a 77% decrease. SEM micrographs ( Figs. 2a 2 and 2b 2 ) show intragranular acicular precipitates. In LSP sample, these fine precipitates cluster at grain centers, forming a distinct precipitation-free zone (PFZ) around grain boundaries. In HSP sample, intragranular acicular nitrides are more dispersed, and the PFZ is significantly narrowed, almost negligible. Dark field (DF) micrographs and EDS maps (Fig. 2 c) reveal a uniform distribution of abundant nanoscale acicular precipitates within the matrix, with strong spatial correlation between Cr and N signals. Discrete dislocations are observed surrounding or penetrating the precipitates (Fig. 2 c 1 ). HRTEM micrographs (Fig. 2 d, Fig. S4 ) display the acicular width and interfacial contours of these precipitates, showing excellent lattice coherence between precipitates and the surrounding matrix. A high density of lattice defects/distortions is observed within the precipitates. Fast Fourier Transform (FFT) patterns from precipitate-matrix interfaces (Fig. 2 e) reveal orientation relationships of [011] M //[001] CrN and [001] M //[011] CrN , confirming the face-centered cubic (fcc) crystal structure of CrN precipitates and their crystallographic orientation with the ferrite matrix, in full agreement with previous reports [ 24 , 25 ]. Figure 3 characterizes the grain boundary microstructure and deformation response of HSP sample. In Fig. 3a 1 and Fig. S5 , ferrite grain boundaries are enveloped by submicron-sized blocky phases and nanoscale banded structures. Electron diffraction patterns ([-111] and [0–11] zone axes, see Figs. 3a 2 and 3a 3 ) confirm these phases as fcc austenite (γ). Elemental mapping (Fig. 3 b) shows nitrogen enrichment in grain boundary austenite, while nickel distribution is uniform across phases. The nitrogen partition coefficient between austenite (~ 8 at.%) and ferrite (~ 3.5 at.%) is approximately 2.5. This value is significantly lower than those typically reported for conventional duplex stainless steels [ 26 ]. Figures 3 c and 3 d show other grain boundary regions. Elongated CrN precipitates with > 20 at.% nitrogen are present along some grain boundaries (see line profile in Fig. S6 ). In regions without distinct precipitates, nitrogen atom clusters (> 5 at.%) are observed at grain boundaries (Fig. 3 d). Overall, these LPBF samples exhibit grain boundaries dominated by submicron-sized blocky austenite, with coexisting minor nitride precipitates or nitrogen segregation. Figure 3 e and Fig. S7 illustrates the mechanical stability of grain boundary austenite near the tensile fracture. Deformed austenite contains numerous elongated acicular structures. HRTEM micrographs (R1–R4) and corresponding FFT patterns confirm these as lamellar twins [ 27 ]. The FFT pattern of region R2 and R3 shows additional diffraction spots (1/3, 2/3 ), indicating local defects at the twin boundary interface, identified as a long-period ordered hexagonal close-packed (hcp) 9R structure [ 28 , 29 ]. Solid-liquid interface velocity dominating nitrogen supersaturation Coupled thermodynamic, kinetic, and multiphysics simulations were performed to analyze phase stability, nitrogen redistribution, and precipitation kinetics under non-equilibrium LPBF conditions, establishing correlations between laser scanning parameters and thermal conditions. Simulations of the Fe-Cr-Ni-Mo-Si-N system ( Fig. S8a ) predicted primary BCC ferrite solidification at ~ 1450°C, followed by FCC austenite formation at ~ 1365°C for 0.63 wt.% N, consistent with observed microstructures ( Figs. 2a 2 , 2b 2 and 3a 1 ). Dictra simulations (Fig. 4 a) investigate nitrogen redistribution kinetics across a ferrite (20 µm)-austenite (2 µm) diffusion couple under various cooling rates. High cooling rates (≥ 2600°C/s) suppress nitrogen diffusion, leading to steep concentration gradients and higher retained nitrogen in ferrite grain interiors. Specifically, at a cooling rate of 10 4 o C/s, high nitrogen concentration was maintained up to 10 µm from ferrite grain centers before declining, with significant nitrogen drops near ferrite grain boundaries supporting experimental PFZ in LSP ( Figs. 2a 2 ). Grain refinement in HSP (< 10 µm vs. ~26 µm in LSP, Fig. S2 ), combined with extreme cooling, inhibited PFZ formation, supporting in a more uniform distribution of nitride precipitates in HSP ( Fig. 2b 2 ). This kinetic trapping was quantified ( Fig. S8b ) at 10 4 o C/s, nitrogen concentration remained over three times the equilibrium value at 1100°C, demonstrating severe supersaturation. This high supersaturation drove prolific, homogeneous nanoscale CrN nucleation. Nitride growth modulated chemistry, as Cr depletion zones formed ( Fig. S8c ) with widths dependent on precipitate size. Specifically, increasing lamella thickness from tens to hundreds of nanometers significantly expanded the adjacent Cr-depleted zone from less than 0.2 µm to over 1 µm. The above calculation results are highly consistent with prior work by N. Pettersson et al. [ 30 ] and S. Hertzman et al. [ 31 ] on compositionally analogous steels. Simulated 2D cross-sectional views of temperature fields and melt pool morphology (Figs. 4 c, 4 d) for LSP and HSP conditions reveal that increased SP elongates the melt pool, constricting its width and depth (e.g., width from ~ 108 µm to 81 µm). The temperature profile along the scanning direction (Fig. 4 e) quantified the in-plane thermal gradient ( G x ), which decrease with increasing SP. Thermal cycles (Fig. 4 f) derived average cooling rates of ~ 1.5×10 6 °C/s for LSP and ~ 2.8×10 6 °C/s for HSP. This increase in cooling rate, despite decreasing G x , is resolved by the relationship R ≈ G × v , where v is the solidification front velocity scaling with SP [ 32 ]. The ~ 2.3-fold increase in v dominated, increasing the net cooling rate ( R ) by ~ 1.8-fold. Thus, the kinematic component ( v ) overwhelmingly governs cooling intensity under extreme LPBF conditions. Breaking strength-corrosion trade off under LPBF process Figure 5 illustrates the synergistic mechanical and electrochemical properties achieved in the LSP and HSP samples. As shown in Fig. 5 a, both HSP and LSP samples exhibit ultra-high tensile strengths, reaching 1700 MPa and 1300 MPa, respectively. This performance surpasses that of other advanced steel systems, such as NiAl steel and WA-DED 430 (Fig. 5 b), highlighting the significant advantages conferred by the fine dense CrN lamellae strengthening in our LPBF samples [ 33 – 35 ]. These results demonstrate that SP serves as a critical parameter for tailoring the mechanical properties, enabling targeted optimization of strength and ductility. Furthermore, the potentiodynamic polarization curves (Fig. 5 c) and the quantitative comparison of pitting potential ( E pit ) versus tensile strength (Fig. 5 d) confirm the stable and excellent corrosion resistance of both HSP and LSP samples. Notably, the pitting potential remains consistently high at approximately 1000 mV, unaffected by variations in SP. This exceptional pitting resistance, coupled with ultra-high tensile strength, positions our material significantly above all other reference materials, including DSSs, In718, C276 and CoCrFeNiTi high entropy alloys [ 36 , 37 ]. This synergistic combination of properties unequivocally indicates that the HSP steel achieves a breakthrough in simultaneously enhancing both strength and corrosion resistance. In summary, this study conclusively demonstrates that interface velocity-driven non-equilibrium nitrogen supersaturation during additive manufacturing provides an efficient and universal strategy to overcome the long-standing strength-corrosion trade-off in metallic alloys. The engineered non-equilibrium microstructure, characterized by ultra-high tensile strength (~ 1700 MPa) and exceptional pitting corrosion resistance (~ 1000 mV), decisively breaks the classical strength-corrosion trade-off. This work offers a powerful new pathway for designing hierarchical microstructures and accelerating the rational development of next-generation high-performance alloys through additive manufacturing. Discussion Contribution of non-equilibrium precipitation for strength increase To quantify strengthening mechanisms, we first establish a baseline for LPBF-processed high nitrogen ferritic stainless steels without extensive precipitation hardening. Nie et al. [ 38 ] reported a ~ 800 MPa yield strength in LPBF-fabricated duplex stainless steel (predominantly ferritic, ~ 96%), attributed to grain refinement, solid solution strengthening, and high geometrically necessary dislocation (GND) density (~ 1.4 × 10 14 m -2 ). He et al. [ 3 ] corroborated this, showing dislocation density saturates at 10 13 -10 14 m -2 for scan speeds ≥ 700 mm/s, with minimal laser power impact. This consistent high dislocation density ( Fig. 2c 1 ) is a fundamental characteristic of rapid solidification LPBF, forming a reliable baseline strength. Assuming similar high dislocation densities (~ 10 14 m⁻²) for our LSP and HSP conditions, the dislocation strengthening ( Δσ dis .) is comparable. Using the Taylor hardening model [ 39 ] with ferritic parameters (Taylor factor M = 2.7, shear modulus G = 83 GPa, Burgers vector b = 0.25 nm, and obstacle efficiency α = 0.475), Δσ dis. is estimated at 315 MPa. For solid solution strengthening (Δσ SS ), the empirical relationship Δσ SS = 1103.45C + 1103.45N + 25.8Si + 19.2Ni + 16.9Mn + 15.9Mo + 2.6Cr (in MPa, with element concentrations in wt. %) [ 40 , 41 ] was applied. Adopting a realistic nitrogen content of 0.1 wt.% ( Fig. S8a ), due to its low equilibrium solubility in bcc-Fe, Δσ SS totals 439 MPa. Grain refinement strengthening (Δσ hp ) was quantified via the Hall-Petch relationship [ 42 , 43 ]. With a Hall-Petch coefficient (k y ) of 0.55 MPa·m 1/2 [ 44 , 45 ] and measured average grain sizes of 26.1 µm (LSP) and 6.2 µm (HSP), Δσ hp values are approximately 108 MPa and 220 MPa, respectively. Table S3 summarizes the quantitative contributions to yield strength. This highlights that non-equilibrium CrN precipitation, activated at higher scan speed, is the dominant mechanism, accounting for ~ 72% of the total strength increment and critically contributing to superior mechanical properties for HSP sample. Interface velocity-driven interstitial solute trapping accounting for counterintuitive enhanced precipitation Conventional wisdom dictates that slower cooling rates favor precipitation by allowing sufficient time for solute diffusion and phase formation, a principle well-established in phase transformation kinetics and observed in traditional alloy processing [ 46 – 48 ]. Similarly, in additive manufacturing, lower scan speeds (implying slower cooling) often facilitate detrimental precipitate formation [ 5 , 49 – 51 ]. However, this study reveals a counterintuitive finding: increasing laser scan speed significantly enhances nanoscale CrN precipitation, as evidenced by the remarkably higher density of nanoscale precipitates in the HSP sample compared to the LSP (Fig. 2 ) and its contribution to strength (see above subsection). This "higher scan speed, more CrN precipitation" phenomenon (illustrated in Fig. S9 ) reveals substantial nitrogen enrichment in the ferritic matrix, significant grain refinement, dramatically increased density of intragranular CrN nanoprecipitates, and suppressed PFZ under high scan speed, while maintaining a continuous corrosion-resistant grain boundary austenite network. This anomalous precipitation behavior is driven by the unique non-equilibrium solidification and thermal cycles inherent to LPBF (Figs. 4 , S8). The fundamental mechanism involves two key aspects. Firstly, accelerated solidification, primarily controlled by the solid-liquid interface velocity and regulated by laser scan speed, leads to "solid-liquid interface velocity dominating nitrogen supersaturation" as supported by our simulations (Figs. 4 c, 4 d). Secondly, accelerated solidification dramatically suppresses long-range diffusion of interstitial nitrogen, kinetically trapping a vastly higher concentration of solutes within the ferritic matrix, forming a highly supersaturated solid solution. The resultant enormous chemical driving force (Fig. 4 b) then promotes the prolific, homogeneous nucleation of fine, coherent metastable CrN precipitates, favored over Cr 2 N due to its lower coherent interfacial energy with the ferrite matrix ( Fig. 2d 1 ). These fine precipitates contribute substantially to the ~ 600 MPa strength increase (Table S3 ) in the HSP sample. This is further facilitated by the rapid thermal cycles of LPBF, which generate substantial thermal stress and plastic strain, leading to a high density of dislocations (Fig. 2 c) and vacancies [ 52 – 54 ] that act as preferential nucleation sites, thereby enhancing precipitation nucleation. Therefore, this counterintuitive enhanced precipitation is a synergistic outcome of interface velocity-driven ultra-fast solidification, which simultaneously creates (i) high supersaturation (thermodynamic driving force) and (ii) high defect density (kinetic pathway), collectively facilitating a burst of fine-scale precipitation otherwise kinetically inaccessible. Ultra-fast solidification induced nitrogen-deficient austenite and associated toughening via twinning Commercial 2507/3207 high-nitrogen duplex stainless steels exhibit extreme austenite stability, precluding deformation-induced twinning or martensitic transformation [ 48 , 55 ]. Our LPBF process overcomes this by engineering a metastable grain boundary austenite, activating additional energy-dissipating deformation mechanisms. LPBF's inherent non-equilibrium solidification is key. Extreme solidification rates kinetically trap nitrogen in the ferritic matrix, suppressing its partitioning to austenite. Consequently, grain boundary austenite is significantly nitrogen-depleted (~ 1 wt.% N, Fig. 4 a) compared to conventional counterparts. This compositional shift is critical, as nitrogen potently modulates austenite stacking fault energy (SFE) [ 56 ]. While not monotonic, DFT calculations [ 57 ] suggest that increasing nitrogen to 1.0 wt.% promotes twinning over cross-slip, consistent with our observations (Fig. 3 e). This nitrogen-deficient state substantially lowers SFE, rendering the boundary austenite mechanically metastable. Beyond deformation twins, a long-range ordered 9R structure was observed at twin boundaries ( Fig. 3e 3 ). Both nanoscale twins and 9R structures act as potent obstacles to dislocation motion, enhancing work hardening (TWIP effect) [ 29 , 58 ]. Simultaneously, grain boundary austenite impedes crack propagation via ductile bridging and blunting [ 59 ]. Thus, toughening arises from a process-enabled synergy: LPBF-induced N-depletion creates low-SFE metastable austenite, activating stacking faulting and deformation twinning. This, combined with austenite's metastable characteristics and enhanced CrN precipitation, comprehensively explains the excellent strength-ductility synergy under HSP condition. Our work highlights that in AM processes under extreme non-equilibrium, traditional precipitation kinetics can be inverted, enabling novel microstructural design via precise thermal kinetic control. Non-equilibrium solidification engineering grain boundary nitrogen for boosted corrosion resistance​ This study reveals exceptional pitting corrosion resistance (Epit ≈ 1000 mV) HSP sample, despite a high density of chromium-rich nitride precipitates. This contradicts the established paradigm where CrN precipitation typically degrades localized corrosion resistance due to adjacent chromium depletion [ 60 , 61 ]. The anomaly stems from a unique, multi-scale corrosion protection architecture engineered by LPBF. First, at grain boundaries (GBs), the refined microstructure avoids vulnerable corrosion pathways. TEM analysis shows the GBs are either enriched with nitrogen-stabilized austenite or decorated with nanoscale, continuous Cr, N-rich layers (Figs. 3 , S5 and S6). This configuration, particularly Cr, N enrichment in austenite, results in a localized pitting resistance equivalent (PREN = %Cr + 3.3 × %Mo + 16 × %N) [ 62 ] significantly higher than the ferritic grain interior. Consequently, these GBs act as corrosion-resistant barriers, suppressing intergranular attack and hindering pit propagation. Second, within ferritic grains, the high-density nanoprecipitates differ fundamentally from coarse, equilibrium nitrides. Formed under extreme non-equilibrium solidification, these precipitates are ultra-fine and exhibit coherent/semi-coherent interfaces ( Fig. 2d 2 ). Crucially, the surrounding chromium-depleted zone is extremely narrow and chemically mild due to suppressed diffusion and non-equilibrium composition ( Fig. S8c ). The adjacent matrix's chromium concentration remains sufficiently high to maintain passivity [ 7 , 63 ]. Thus, these intragranular nanoprecipitates, while providing potent Orowan strengthening, do not create the extensive, interconnected chromium-depleted networks that initiate stable pitting, unlike coarse counterparts in conventionally processed materials. In summary, high corrosion resistance is preserved through a synergistic, multi-level defense: (1) corrosion-resistant GBs with high localized PREN, and (2) intragranular nanoprecipitates with minimal detrimental impact on the local electrochemical environment. This unique microstructural configuration, enabled by rapid solidification and defect-assisted precipitation kinetics of high-speed LPBF, successfully decouples nitrides' strengthening effect from their traditionally deleterious impact on corrosion resistance. Methods Powder preparation & LPBF processing The high nitrogen stainless steel powder for LPBF printing was prepared by vacuum induction melting gas atomization (VIGA) method, which was finished in the Institute of Metal Research, Chinese Academy of Sciences. The chemical compositions of the powder were determined by various methods (ICP method, ONH and CS analyzers) as shown in Table S1 . This composition is characterized by a high nitrogen content, approximately 0.6 wt.%, hence the LPBF steel is conveniently named 6N. After powder preparation, the powder undergoes further sieving to select powder with a particle size of approximately 15–53 µm for printing. The powders were dried for at least 5 h in a vacuum oven at 70 ℃ before additive manufacturing fabrication. The LPBF processing was performed via a 3D additive manufacturing system (Hans-M-100, Hans Laser) equipped with a continuous Ytterbium fiber laser (wavelength: 1064 nm, spot diameter: 50 µm). Numerous orthogonal experiments (hatch spacing: 80–120 µm, scanning speed: 200–1500 mm/s, laser power 200–300 W, layer thickness: 30–60 µm) have been explored previously, and a relatively optimal printing parameter space with varying scanning speed has been obtained, as presented in Table S2 . Apart from the scanning speed, the printing power ( P ), layer thickness ( t ), hatch spacing ( H ), and printing strategy remain unchanged at P = 250 W, t = 50 µm, H = 80 µm, respectively. The commonly bidirectional scanning strategy with a rotation angle of 67 o between the upper and lower layers was used for minimizing texture and anisotropy as shown in Fig. S1a . All samples were fabricated in a high-purity flowing argon atmosphere for avoiding oxidation or nitriding. A 316L-type stainless steel substrate (110 mm × 110 mm × 13 mm) was used due to its compatibility with stainless steel materials. Plate tensile & potentiodynamic polarization tests These cuboid samples with different scanning parameters were cut from the 316L substrate and were further cut into several plate tensile samples using wire electrical discharge machining. The plane of the tensile sample was parallel to the substrate and other sampling details like sample size and thickness are shown in Fig. S1b . All tensile samples were mechanically ground using SiC sandpapers ranging from 180 to 2000 grit to obtain flat and clean surfaces. Tensile tests were conducted on a universal testing machine equipped with an optical extensometer. The tensile strain rate was 10 − 3 s − 1 . After tensile testing, the clamping planes of the fractured samples were used for the potentiodynamic polarization experiment. All samples were mechanically re-polished using silicon carbide paper with grit sizes ranging from 180 to 2000 and finished by using 2.5 µm diamond suspension for 30 minutes and vibratory polishing for 2 hours. The electrochemical measurements were conducted on an electrochemical workstation (Versa STAT 3 F) using a three-electrode cell. The working electrode was an investigated sample that was fixed in a holder with an exposing area of ~ 0.28 cm 2 , a platinum plate was used as the counter electrode, and a saturated Ag/AgCl electrode was used as the reference electrode. Potentiodynamic polarization tests were carried out in 3.5% NaCl electrolyte, with a standard scan rate of 0.167 mV/s, an initial potential of − 0.6 V, and a final potential of 1.2 V against the open circuit potential (OCP). The potentiodynamic polarization test and electrochemical resistance spectroscopy (EIS) were conducted after OCP by immersing the sample for 30 min in the electrolyte to wait for the system to stabilize. EIS measurements were performed by using a frequency range from 10 5 Hz to 10 − 2 Hz, and the root mean square (RMS) potential amplitude is 10 mV. The electrochemical measurements were conducted at room temperature (22–24°C). Microstructural characterization The morphologies and crystal orientation of these 6N samples were observed by a field emission scanning electron microscope (SEM, Apreo 2, FEI) equipped with a high-speed electron backscatter diffraction (EBSD, Velocity, EDAX) detector and a scanning transmission electron microscope (STEM) detector. EBSD maps were obtained with varying step sizes of 0.05 and 0.1 µm according to collection region areas. Raw EBSD data were analyzed using the OIM (EDAX) software. For TEM sampling, both FIB and electrochemical jet polishing methods were used. Some 3 mm disks were twin-jet electro-polished by a TenuPol-5 (Struers) in an 8% HClO 4 ethanol solution at 25°C and 20 V. The TEM lamellas were prepared via FIB (Helios 600i, FEI) with Pt protection, coarse milling at 30 kV, and final polishing at 5 kV to ensure electron transparency. A 200 kV transmission electron microscope (TEM, Talos F200X G2, FEI) equipped with high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) functionality, and energy-dispersive X-ray spectrometer (EDS) was employed for microstructural analysis. Simulation methods A three-dimensional thermo-fluidic simulation of the single-track melt-pool was performed. The model solves the governing equations of mass, momentum, and energy conservation to predict the transient temperature field, melt-pool geometry, and thermal history under different laser scan speeds. A Gaussian volumetric heat source was employed to simulate the laser energy input. Temperature-dependent thermophysical properties (e.g., thermal conductivity, specific heat, viscosity) of the alloy were incorporated from the Thermo-Calc FEDEMO V5 database. The simulations were conducted for different scanning speed, corresponding to the experimental conditions. Thermodynamic and kinetic calculations were employed to understand the phase stability and solute redistribution behavior under the extreme non-equilibrium solidification conditions of LPBF. Equilibrium phase diagram calculations and driving force analyses were performed using the Thermo-Calc software with a relevant thermodynamic database (TCFE7). A vertical section of the multi-component system (Fe-Cr-Ni-Mo-Si-N) was calculated to map the stability regions of ferrite (BCC) and austenite (FCC) as a function of nitrogen content and temperature. Furthermore, the normalized driving forces for the precipitation of Cr 2 N and CrN phases from the ferrite matrix were calculated based on its nominal composition. Diffusion simulations were performed using the Dictra module. A one-dimensional diffusion couple model was constructed, consisting of ferrite (BCC_A2), Cr 2 N (HCP) and austenite (FCC_A1) phases. Simulations modeled the continuous cooling from 1360 o C at rates representative of the melt pool simulation results. Declarations Data availability Data will be made available on request. Declaration of competing interest The authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper. Acknowledgements M. Yang acknowledges the support from the Guangdong Basic and Applied Basic Research Foundation (Grant No. 2023A1515240048), the Shenzhen Science and Technology Program (Grant No. JCYJ20240813094806009) and the National Natural Science Foundation of China (Grant No. 52301153); J. Yi acknowledges the support from the National Natural Science Foundation of China (Grant No. 52301154) and the Shenzhen Science and Technology Program (Grant No. JCYJ20220530114400001); M. He acknowledges the support from the Guangdong Basic and Applied Basic Research Foundation (Grant No. 2023A1515110595); Z. Yao acknowledges the support from the National Natural Science Foundation of China (Grant No. 52301042) and the Natural Science Foundation of Top Talent of SZTU (grant No. GDRC202532). All authors gratefully acknowledge the assistance from Dr. Y. Qiu and Dr. D. He at SUStech Core Research Facilities and the electron microscope center of KAIPLE Co. 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Zhu, Long-period stacking ordering induced ductility of nanolamellar TiAl alloy at elevated temperature, Materials Research Letters 11(6) (2023) 414–421. D. Raabe, S. Sandlöbes, J. Millán, D. Ponge, H. Assadi, M. Herbig, P.-P. Choi, Segregation engineering enables nanoscale martensite to austenite phase transformation at grain boundaries: a pathway to ductile martensite, Acta Materialia 61(16) (2013) 6132–6152. K. Krishna Kumar, J. Anburaj, R. Dhanasekar, T. Satishkumar, J. Abuthakir, P. Manikandan, R. Subramanian, Kinetics of Cr2N Precipitation and Its Effect on Pitting Corrosion of Nickel-Free High-Nitrogen Austenitic Stainless Steel, Journal of Materials Engineering and Performance 29(9) (2020) 6044–6052. H. Ha, H. Kwon, Effects of Cr2N on the pitting corrosion of high nitrogen stainless steels, Electrochimica Acta 52(5) (2007) 2175–2180. L. Garfias-Mesias, J. Sykes, C. Tuck, The effect of phase compositions on the pitting corrosion of 25 Cr duplex stainless steel in chloride solutions, Corrosion science 38(8) (1996) 1319–1330. E. Bettini, U. Kivisäkk, C. Leygraf, J. Pan, Study of corrosion behavior of a 2507 super duplex stainless steel: influence of quenched-in and isothermal nitrides, International Journal of Electrochemical Science 9(1) (2014) 61–80. Additional Declarations There is NO Competing Interest. Supplementary Files Supplementarymaterials.docx Dataset 1 Cite Share Download PDF Status: Posted Version 1 posted You are reading this latest preprint version Research Square lets you share your work early, gain feedback from the community, and start making changes to your manuscript prior to peer review in a journal. As a division of Research Square Company, we’re committed to making research communication faster, fairer, and more useful. We do this by developing innovative software and high quality services for the global research community. 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Also discoverable on Platform About Our Team In Review Editorial Policies Advisory Board Help Center Resources Author Services Accessibility API Access RSS feed Manage Cookie Preferences © Research Square 2026 | ISSN 2693-5015 (online) Privacy Policy Terms of Service Do Not Sell My Personal Information {"props":{"pageProps":{"initialData":{"identity":"rs-9249060","acceptedTermsAndConditions":true,"allowDirectSubmit":true,"archivedVersions":[],"articleType":"Article","associatedPublications":[],"authors":[{"id":614825986,"identity":"5d86ba7d-4d2e-4dfb-b853-d47a933c9400","order_by":0,"name":"Mujin 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02:10:17","currentVersionCode":1,"declarations":"","doi":"10.21203/rs.3.rs-9249060/v1","doiUrl":"https://doi.org/10.21203/rs.3.rs-9249060/v1","draftVersion":[],"editorialEvents":[],"editorialNote":"","failedWorkflow":false,"files":[{"id":106473091,"identity":"cf748b59-7c0b-4f18-90c2-8ce278747d48","added_by":"auto","created_at":"2026-04-09 02:21:03","extension":"png","order_by":1,"title":"Figure 1","display":"","copyAsset":false,"role":"figure","size":229428,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eConceptual schematic of nitrogen saturation and microstructural landscape: \u003c/strong\u003e(a) Traditional casting shows quasi-equilibrium solidification, leading to limited, coarse nitrides. (b) In the AM melt pool, scanning speed modulates solidification kinetics: (b\u003csub\u003e1\u003c/sub\u003e) Low scanning speed (LSP) results in metastable austenite, fine nitrides, and precipitation-free zone (PFZ). (b\u003csub\u003e2\u003c/sub\u003e) High scanning speed (HSP) further suppresses atomic diffusion, enabling finer nitrides without PFZ. Note that: Nitrogen redistribution and grain boundary austenite form near the melting temperature (T=T\u003csub\u003em\u003c/sub\u003e), while nitride precipitation occurs at cooling process at lower temperature.\u003c/p\u003e","description":"","filename":"1.png","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/543748d22106aecf10b02c51.png"},{"id":106473092,"identity":"798e30ff-e231-41fd-889f-929f5a8a0a3c","added_by":"auto","created_at":"2026-04-09 02:21:03","extension":"png","order_by":2,"title":"Figure 2","display":"","copyAsset":false,"role":"figure","size":1082586,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMultiscale microstructural characterization of grains and acicular precipitates in LSP and HSP samples: \u003c/strong\u003e(a, b) IPF and SEM micrographs. (c) DF micrograph of acicular CrN precipitates and EDS maps (Cr, N). (d) HRTEM micrographs of CrN and ferrite matrix. (e, f) FFT patterns showing the CrN/matrix orientation relationship.\u003c/p\u003e","description":"","filename":"2.png","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/74f5322afbf50f42dcbce687.png"},{"id":106473093,"identity":"3ef5b88c-2ce9-4192-9bbb-9889863121b8","added_by":"auto","created_at":"2026-04-09 02:21:03","extension":"png","order_by":3,"title":"Figure 3","display":"","copyAsset":false,"role":"figure","size":982311,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMultiscale microstructural characterization of GB in HSP sample:\u003c/strong\u003e (a) Morphology of GB austenite with corresponding electron diffraction patterns. (b) EDS maps and line scan profiles of Ni and N at the ferrite/austenite interface, from the enlarged region in (a). (c, d) Morphology and nitrogen EDS mapping of two distinct GB characteristics: (c) showing GB nitride precipitation, and (d) showing GB nitrogen segregation. (e) Morphology of deformation-induced twins within metastable GB austenite, observed near the tensile fracture, with corresponding high-resolution micrograph and FFT pattern.\u003c/p\u003e","description":"","filename":"3.png","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/b2b2f99bfe0eac9a072bf6ac.png"},{"id":106473095,"identity":"bf7a5cf9-24fa-4d03-b7d4-99cc0535407c","added_by":"auto","created_at":"2026-04-09 02:21:03","extension":"png","order_by":4,"title":"Figure 4","display":"","copyAsset":false,"role":"figure","size":433293,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eThermodynamic, kinetic and multi-physics simulations: \u003c/strong\u003e(a) Nitrogen concentration profiles across a ferrite (20 μm)-austenite (2 μm) diffusion couple under varying cooling rates; (b) Calculated transformation driving force for the formation of CrN and Cr\u003csub\u003e2\u003c/sub\u003eN as a function of temperature, based on the ferrite composition at 1360 °C. (c, d) 2D cross-sectional views (X-Y, X-Z and Y-Z plane) of temperature field and melt pool morphology at LSP and HSP, respectively. Note that: Simulations used the Thermo-Calc FEDEMO database and simplified Fe-20Cr-8Ni (wt.%) composition. (e) Temperature distribution profiles along the scanning direction at a fixed temperature. (f) Thermal cycles (temperature vs. time) at a fixed point.\u003c/p\u003e","description":"","filename":"4.png","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/d8587d20c3b7e5c6e766664d.png"},{"id":106725344,"identity":"9e13fd70-ac75-4941-aec2-b104b93b0ee4","added_by":"auto","created_at":"2026-04-12 18:32:32","extension":"png","order_by":5,"title":"Figure 5","display":"","copyAsset":false,"role":"figure","size":238964,"visible":true,"origin":"","legend":"\u003cp\u003e\u003cstrong\u003eMechanical and corrosion properties of the LPBF-fabricated LSP and HSP samples. \u003c/strong\u003e(a) Engineering stress-strain curves (b) Ultimate tensile strength vs. uniform elongation for the LSP and HSP samples (red star), compared with additively manufactured steels (AM-S441), wire-arc directed energy deposition steels (WA-DED 430) [33, 34], and others from literature [35]. (c) Potentiodynamic polarization curves tested in a 3.5 wt.% NaCl solution. (d) Pitting potential (E\u003csub\u003epit\u003c/sub\u003e) vs. tensile strength correlated with commercial alloys including duplex stainless steel, IN718, C276, and AM CoCrFeNiTi high-entropy alloy (HEA) [36, 37].\u003c/p\u003e","description":"","filename":"5.png","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/c41a1b2709dee99e6fe31dc3.png"},{"id":106727295,"identity":"c310a187-d5bf-4fd5-ab66-2da89a5dadb4","added_by":"auto","created_at":"2026-04-12 18:38:35","extension":"pdf","order_by":0,"title":"","display":"","copyAsset":false,"role":"manuscript-pdf","size":4050404,"visible":true,"origin":"","legend":"","description":"","filename":"manuscript.pdf","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/0d7d4ed7-51a8-4ab9-8ceb-8b794ba944c3.pdf"},{"id":106473097,"identity":"cd659843-f471-4092-b8fb-0751091505d7","added_by":"auto","created_at":"2026-04-09 02:21:04","extension":"docx","order_by":1,"title":"","display":"","copyAsset":false,"role":"supplement","size":94552568,"visible":true,"origin":"","legend":"Dataset 1","description":"","filename":"Supplementarymaterials.docx","url":"https://assets-eu.researchsquare.com/files/rs-9249060/v1/e2b9158d5b3b4ea3a1afcbaf.docx"}],"financialInterests":"There is \u003cb\u003eNO\u003c/b\u003e Competing Interest.","formattedTitle":"Interface velocity-driven non-equilibrium nitrogen supersaturation in additive manufacture: a universal strategy for breaking strength-corrosion trade off","fulltext":[{"header":"Introduction","content":"\u003cp\u003e \u003cdiv class=\"BlockQuote\"\u003e \u003cp\u003eLaser powder bed fusion (LPBF) has unlocked new frontiers in metallurgy by leveraging its characteristic extreme thermal cycles\u0026mdash;featuring rapid solidification (~\u0026thinsp;10\u003csup\u003e3\u003c/sup\u003e\u0026ndash;10\u003csup\u003e6\u003c/sup\u003e K/s) and repeated remelting\u0026mdash;to create non-equilibrium microstructures and defect landscapes unattainable through conventional processing [\u003cspan citationid=\"CR1\" class=\"CitationRef\"\u003e1\u003c/span\u003e, \u003cspan citationid=\"CR2\" class=\"CitationRef\"\u003e2\u003c/span\u003e]. This ability to dictate non-equilibrium phase transformations presents a compelling opportunity to revisit and potentially solve persistent issues in alloy design, including the control of detrimental secondary phase precipitation in systems like high nitrogen duplex stainless steels (DSSs) [\u003cspan citationid=\"CR3\" class=\"CitationRef\"\u003e3\u003c/span\u003e, \u003cspan citationid=\"CR4\" class=\"CitationRef\"\u003e4\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eHigh nitrogen DSSs with their balanced ferrite-austenite microstructure, are prized for their excellent combination of toughness and corrosion resistance, serving critical roles in marine, chemical, and energy industries [\u003cspan citationid=\"CR5\" class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan citationid=\"CR6\" class=\"CitationRef\"\u003e6\u003c/span\u003e]. However, their performance is intrinsically constrained by a classic trade-off: conventional strengthening methods, such as cold working or precipitation hardening via secondary phases, often degrade their corrosion resistance. A central challenge lies in controlling the precipitation of chromium nitrides (Cr\u003csub\u003e2\u003c/sub\u003eN/CrN) [\u003cspan citationid=\"CR7\" class=\"CitationRef\"\u003e7\u003c/span\u003e, \u003cspan citationid=\"CR8\" class=\"CitationRef\"\u003e8\u003c/span\u003e]. During conventional processing, these nitrides typically precipitate in the temperature range of 700\u0026ndash;1000\u0026deg;C, either during slow cooling or isothermal aging. Even under rapid quenching, so-called \u0026ldquo;quenched-in\u0026rdquo; nitrides can form within the ferrite matrix, as the high cooling rate suppresses nitrogen diffusion into austenite, leading to supersaturation and subsequent intragranular precipitation [\u003cspan citationid=\"CR9\" class=\"CitationRef\"\u003e9\u003c/span\u003e, \u003cspan citationid=\"CR10\" class=\"CitationRef\"\u003e10\u003c/span\u003e]. While refining the nitride morphology, these precipitates still deplete chromium in adjacent areas and are widely reported to impair pitting corrosion resistance. Therefore, avoiding or minimizing nitride formation has been a guiding principle in the thermo-mechanical processing of high nitrogen DSSs.\u003c/p\u003e \u003cp\u003eIn pursuit of this goal, several strategies have been explored to mitigate this trade-off. Conventional approaches often focus on precise thermo-mechanical processing routes to control the nucleation and growth of deleterious phases [\u003cspan citationid=\"CR11\" class=\"CitationRef\"\u003e11\u003c/span\u003e, \u003cspan citationid=\"CR12\" class=\"CitationRef\"\u003e12\u003c/span\u003e]. Alloy design innovations, such as optimizing the Cr/Ni/Mn/N ratio or adding nitride-stabilizing elements like Cu, have also been attempted to refine precipitate distribution [\u003cspan additionalcitationids=\"CR14\" citationid=\"CR13\" class=\"CitationRef\"\u003e13\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR15\" class=\"CitationRef\"\u003e15\u003c/span\u003e]. More recently, advanced manufacturing techniques like directed energy deposition (DED) have been employed to create graded microstructures, yet control over nanoscale precipitation remains challenging [\u003cspan citationid=\"CR16\" class=\"CitationRef\"\u003e16\u003c/span\u003e]. While these methods have met with varying degrees of success, they often involve complex post-processing or still struggle to completely decouple the strength enhancement provided by precipitates from their detrimental effect on corrosion resistance. The fundamental challenge lies in the inherent equilibrium or near-equilibrium nature of these processes, which favors the formation of coarse, chromium-depleting precipitates.\u003c/p\u003e \u003cp\u003eThe LPBF process, with its inherent non-equilibrium conditions, challenges this conventional wisdom. The extreme undercooling and thermal stress inherent to LPBF generate a high density of crystal defects, particularly subgrains, dislocations and vacancies [\u003cspan additionalcitationids=\"CR18 CR19\" citationid=\"CR17\" class=\"CitationRef\"\u003e17\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR20\" class=\"CitationRef\"\u003e20\u003c/span\u003e]. In other alloy systems fabricated by LPBF, such as 316L stainless steel and Inconel 718, these processing-induced defects have been shown to act as potent nucleation sites for secondary phases\u0026mdash;a phenomenon termed defect-assisted precipitation [\u003cspan additionalcitationids=\"CR22\" citationid=\"CR21\" class=\"CitationRef\"\u003e21\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR23\" class=\"CitationRef\"\u003e23\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eThis raises a thought-provoking possibility: Can the unique extreme undercooling and special defect distribution induced by thermal stress in LPBF be leveraged to regulate the non-equilibrium supersaturation of nitrogen in the ferrite phase, guide nitride precipitation into dispersed, fine, and favorable morphologies, thereby enhancing strength without compromising corrosion resistance? Specifically, as a key factor determining the solidification rate, the solid-liquid interface migration rate\u0026mdash;controlled by laser scanning speed\u0026mdash;may serve as an effective means to regulate this supersaturation behavior. However, for high-nitrogen stainless steel processed by LPBF, the influence of laser scanning speed on the microstructure-property relationship remains to be explored.\u003c/p\u003e \u003cp\u003eHerein, we demonstrate that LPBF effectively overcomes the traditional strength-corrosion trade-off in a novel high-nitrogen DSSs. We reveal that interface velocity-driven non-equilibrium nitrogen supersaturation, precisely controlled by laser scan speed, promotes the formation of a high dense coherent CrN lamella and substantial grain refinement, concurrently narrowing the precipitation-free zone. Multi-scale characterization shows these microstructural features synergistically lead to a remarkable\u0026thinsp;~\u0026thinsp;400 MPa increase in yield strength, reaching ultra-high strength level, while maintaining exceptional pitting corrosion resistance (E\u003csub\u003epit\u003c/sub\u003e \u0026asymp; 1000 mV). This work not only provides a new pathway for developing high-performance high-nitrogen stainless steels via additive manufacturing but also offers fundamental insights into phase transformation control under extreme processing conditions.\u003c/p\u003e \u003c/div\u003e \u003c/p\u003e"},{"header":"Results","content":"\u003cdiv id=\"Sec3\" class=\"Section2\"\u003e \u003ch2\u003eNon-equilibrium nitrogen supersaturation strategy in additive manufacturing\u003c/h2\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e conceptually illustrates how LPBF regulates non-equilibrium nitrogen partitioning between ferrite and austenite phases, contrasting it with traditional casting. At a constant nitrogen concentration (0.63 wt.%), traditional casting (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003ea) involves quasi-equilibrium solidification, where nitrogen diffuses to achieve near-equilibrium, resulting in significant nitrogen enrichment in austenite and limited, coarsened nitride precipitation at grain boundary (GB) and other defect sites.\u003c/p\u003e \u003cp\u003eIn contrast, the AM melt pool (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb) operates under laser scanning speed (SP)-modulated non-equilibrium solidification kinetics, characterized by significantly much higher solidification and cooling rates, even at low scanning speed (LSP) condition (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e). This rapid process hinders nitrogen diffusion, leading to a non-equilibrium state where ferrite is nitrogen-supersaturated (see point B as arrow pinpointed in nitrogen concentration legend, Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003e) and austenite is slightly nitrogen-deficient (point B'). During subsequent rapid cooling, CrN nucleates and grows. Under LSP (Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e), relatively large grains (tens of micrometers) form with fine, dispersed CrN lamella, often accompanied by a distinct precipitation-free zone (PFZ) near the austenite/ferrite interface. Increasing scanning speed (HSP in Fig.\u0026nbsp;\u003cspan refid=\"Fig1\" class=\"InternalRef\"\u003e1\u003c/span\u003eb\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e) further elevates solidification rates, inducing significant grain refinement (\u0026lt;\u0026thinsp;10 \u0026micro;m) and leading to finer, more dispersed CrN precipitates. This combined refinement significantly narrows or eliminates the PFZ, representing a key microstructural difference between LSP and HSP conditions.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eRegulating non-equilibrium nitrogen supersaturation via laser scan speed\u003c/h3\u003e\n\u003cp\u003eMicrostructural characterization results in Figs.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003e, \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e, S2 and S3 reveal the impact of SP on nitrogen saturation. IPF maps (\u003cb\u003eFig.\u0026nbsp;2a\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e and \u003cb\u003e2b\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e) show no significant texture differences, but increasing SP significantly refines grains. Grain size statistics (\u003cb\u003eFig. S2g\u003c/b\u003e, \u003cb\u003eS2h\u003c/b\u003e) confirm average grain size reduction from 26\u0026thinsp;\u0026plusmn;\u0026thinsp;6 \u0026micro;m (LSP) to 6\u0026thinsp;\u0026plusmn;\u0026thinsp;3 \u0026micro;m (HSP), a 77% decrease. SEM micrographs (\u003cb\u003eFigs.\u0026nbsp;2a\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e and \u003cb\u003e2b\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e) show intragranular acicular precipitates. In LSP sample, these fine precipitates cluster at grain centers, forming a distinct precipitation-free zone (PFZ) around grain boundaries. In HSP sample, intragranular acicular nitrides are more dispersed, and the PFZ is significantly narrowed, almost negligible.\u003c/p\u003e \u003cp\u003eDark field (DF) micrographs and EDS maps (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec) reveal a uniform distribution of abundant nanoscale acicular precipitates within the matrix, with strong spatial correlation between Cr and N signals. Discrete dislocations are observed surrounding or penetrating the precipitates (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ec\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e). HRTEM micrographs (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ed, \u003cb\u003eFig. S4\u003c/b\u003e) display the acicular width and interfacial contours of these precipitates, showing excellent lattice coherence between precipitates and the surrounding matrix. A high density of lattice defects/distortions is observed within the precipitates. Fast Fourier Transform (FFT) patterns from precipitate-matrix interfaces (Fig.\u0026nbsp;\u003cspan refid=\"Fig2\" class=\"InternalRef\"\u003e2\u003c/span\u003ee) reveal orientation relationships of [011]\u003csub\u003eM\u003c/sub\u003e//[001]\u003csub\u003eCrN\u003c/sub\u003e and [001]\u003csub\u003eM\u003c/sub\u003e//[011]\u003csub\u003eCrN\u003c/sub\u003e, confirming the face-centered cubic (fcc) crystal structure of CrN precipitates and their crystallographic orientation with the ferrite matrix, in full agreement with previous reports [\u003cspan citationid=\"CR24\" class=\"CitationRef\"\u003e24\u003c/span\u003e, \u003cspan citationid=\"CR25\" class=\"CitationRef\"\u003e25\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e \u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003e characterizes the grain boundary microstructure and deformation response of HSP sample. In \u003cb\u003eFig.\u0026nbsp;3a\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e and \u003cb\u003eFig. S5\u003c/b\u003e, ferrite grain boundaries are enveloped by submicron-sized blocky phases and nanoscale banded structures. Electron diffraction patterns ([-111] and [0\u0026ndash;11] zone axes, see \u003cb\u003eFigs.\u0026nbsp;3a\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e and \u003cb\u003e3a\u003c/b\u003e\u003csub\u003e\u003cb\u003e3\u003c/b\u003e\u003c/sub\u003e) confirm these phases as fcc austenite (γ). Elemental mapping (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003eb) shows nitrogen enrichment in grain boundary austenite, while nickel distribution is uniform across phases. The nitrogen partition coefficient between austenite (~\u0026thinsp;8 at.%) and ferrite (~\u0026thinsp;3.5 at.%) is approximately 2.5. This value is significantly lower than those typically reported for conventional duplex stainless steels [\u003cspan citationid=\"CR26\" class=\"CitationRef\"\u003e26\u003c/span\u003e].\u003c/p\u003e \u003cp\u003eFigures\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ec and \u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed show other grain boundary regions. Elongated CrN precipitates with \u0026gt;\u0026thinsp;20 at.% nitrogen are present along some grain boundaries (see line profile in \u003cb\u003eFig. S6\u003c/b\u003e). In regions without distinct precipitates, nitrogen atom clusters (\u0026gt;\u0026thinsp;5 at.%) are observed at grain boundaries (Fig.\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ed). Overall, these LPBF samples exhibit grain boundaries dominated by submicron-sized blocky austenite, with coexisting minor nitride precipitates or nitrogen segregation. Figure\u0026nbsp;\u003cspan refid=\"Fig3\" class=\"InternalRef\"\u003e3\u003c/span\u003ee and \u003cb\u003eFig. S7\u003c/b\u003e illustrates the mechanical stability of grain boundary austenite near the tensile fracture. Deformed austenite contains numerous elongated acicular structures. HRTEM micrographs (R1\u0026ndash;R4) and corresponding FFT patterns confirm these as lamellar twins [\u003cspan citationid=\"CR27\" class=\"CitationRef\"\u003e27\u003c/span\u003e]. The FFT pattern of region R2 and R3 shows additional diffraction spots (1/3, 2/3\u0026thinsp;\u0026lt;\u0026thinsp;111\u0026gt;), indicating local defects at the twin boundary interface, identified as a long-period ordered hexagonal close-packed (hcp) 9R structure [\u003cspan citationid=\"CR28\" class=\"CitationRef\"\u003e28\u003c/span\u003e, \u003cspan citationid=\"CR29\" class=\"CitationRef\"\u003e29\u003c/span\u003e].\u003c/p\u003e \u003cp\u003e \u003c/p\u003e\n\u003ch3\u003eSolid-liquid interface velocity dominating nitrogen supersaturation\u003c/h3\u003e\n\u003cp\u003eCoupled thermodynamic, kinetic, and multiphysics simulations were performed to analyze phase stability, nitrogen redistribution, and precipitation kinetics under non-equilibrium LPBF conditions, establishing correlations between laser scanning parameters and thermal conditions. Simulations of the Fe-Cr-Ni-Mo-Si-N system (\u003cb\u003eFig. S8a\u003c/b\u003e) predicted primary BCC ferrite solidification at ~\u0026thinsp;1450\u0026deg;C, followed by FCC austenite formation at ~\u0026thinsp;1365\u0026deg;C for 0.63 wt.% N, consistent with observed microstructures (\u003cb\u003eFigs.\u0026nbsp;2a\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e, \u003cb\u003e2b\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e and \u003cb\u003e3a\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e). Dictra simulations (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ea) investigate nitrogen redistribution kinetics across a ferrite (20 \u0026micro;m)-austenite (2 \u0026micro;m) diffusion couple under various cooling rates. High cooling rates (\u0026ge;\u0026thinsp;2600\u0026deg;C/s) suppress nitrogen diffusion, leading to steep concentration gradients and higher retained nitrogen in ferrite grain interiors. Specifically, at a cooling rate of 10\u003csup\u003e4 o\u003c/sup\u003eC/s, high nitrogen concentration was maintained up to 10 \u0026micro;m from ferrite grain centers before declining, with significant nitrogen drops near ferrite grain boundaries supporting experimental PFZ in LSP (\u003cb\u003eFigs.\u0026nbsp;2a\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e). Grain refinement in HSP (\u0026lt;\u0026thinsp;10 \u0026micro;m vs. ~26 \u0026micro;m in LSP, \u003cb\u003eFig. S2\u003c/b\u003e), combined with extreme cooling, inhibited PFZ formation, supporting in a more uniform distribution of nitride precipitates in HSP (\u003cb\u003eFig.\u0026nbsp;2b\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e). This kinetic trapping was quantified (\u003cb\u003eFig. S8b\u003c/b\u003e) at 10\u003csup\u003e4 o\u003c/sup\u003eC/s, nitrogen concentration remained over three times the equilibrium value at 1100\u0026deg;C, demonstrating severe supersaturation. This high supersaturation drove prolific, homogeneous nanoscale CrN nucleation. Nitride growth modulated chemistry, as Cr depletion zones formed (\u003cb\u003eFig. S8c\u003c/b\u003e) with widths dependent on precipitate size. Specifically, increasing lamella thickness from tens to hundreds of nanometers significantly expanded the adjacent Cr-depleted zone from less than 0.2 \u0026micro;m to over 1 \u0026micro;m. The above calculation results are highly consistent with prior work by N. Pettersson et al. [\u003cspan citationid=\"CR30\" class=\"CitationRef\"\u003e30\u003c/span\u003e] and S. Hertzman et al. [\u003cspan citationid=\"CR31\" class=\"CitationRef\"\u003e31\u003c/span\u003e] on compositionally analogous steels.\u003c/p\u003e \u003cp\u003eSimulated 2D cross-sectional views of temperature fields and melt pool morphology (Figs.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ec, \u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ed) for LSP and HSP conditions reveal that increased SP elongates the melt pool, constricting its width and depth (e.g., width from ~\u0026thinsp;108 \u0026micro;m to 81 \u0026micro;m). The temperature profile along the scanning direction (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ee) quantified the in-plane thermal gradient (\u003cb\u003eG\u003c/b\u003e\u003csub\u003e\u003cb\u003ex\u003c/b\u003e\u003c/sub\u003e), which decrease with increasing SP. Thermal cycles (Fig.\u0026nbsp;\u003cspan refid=\"Fig4\" class=\"InternalRef\"\u003e4\u003c/span\u003ef) derived average cooling rates of ~\u0026thinsp;1.5\u0026times;10\u003csup\u003e6\u003c/sup\u003e \u0026deg;C/s for LSP and ~\u0026thinsp;2.8\u0026times;10\u003csup\u003e6\u003c/sup\u003e \u0026deg;C/s for HSP. This increase in cooling rate, despite decreasing \u003cb\u003eG\u003c/b\u003e\u003csub\u003e\u003cb\u003ex\u003c/b\u003e\u003c/sub\u003e, is resolved by the relationship \u003cb\u003eR\u003c/b\u003e\u0026thinsp;\u0026asymp;\u0026thinsp;\u003cb\u003eG\u003c/b\u003e \u0026times; \u003cb\u003ev\u003c/b\u003e, where \u003cb\u003ev\u003c/b\u003e is the solidification front velocity scaling with SP [\u003cspan citationid=\"CR32\" class=\"CitationRef\"\u003e32\u003c/span\u003e]. The ~\u0026thinsp;2.3-fold increase in \u003cb\u003ev\u003c/b\u003e dominated, increasing the net cooling rate (\u003cb\u003eR\u003c/b\u003e) by ~\u0026thinsp;1.8-fold. Thus, the kinematic component (\u003cb\u003ev\u003c/b\u003e) overwhelmingly governs cooling intensity under extreme LPBF conditions.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e\n\u003ch3\u003eBreaking strength-corrosion trade off under LPBF process\u003c/h3\u003e\n\u003cp\u003eFigure\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003e illustrates the synergistic mechanical and electrochemical properties achieved in the LSP and HSP samples. As shown in Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ea, both HSP and LSP samples exhibit ultra-high tensile strengths, reaching 1700 MPa and 1300 MPa, respectively. This performance surpasses that of other advanced steel systems, such as NiAl steel and WA-DED 430 (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003eb), highlighting the significant advantages conferred by the fine dense CrN lamellae strengthening in our LPBF samples [\u003cspan additionalcitationids=\"CR34\" citationid=\"CR33\" class=\"CitationRef\"\u003e33\u003c/span\u003e\u0026ndash;\u003cspan citationid=\"CR35\" class=\"CitationRef\"\u003e35\u003c/span\u003e]. These results demonstrate that SP serves as a critical parameter for tailoring the mechanical properties, enabling targeted optimization of strength and ductility. Furthermore, the potentiodynamic polarization curves (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ec) and the quantitative comparison of pitting potential (\u003cb\u003eE\u003c/b\u003e\u003csub\u003e\u003cb\u003epit\u003c/b\u003e\u003c/sub\u003e) versus tensile strength (Fig.\u0026nbsp;\u003cspan refid=\"Fig5\" class=\"InternalRef\"\u003e5\u003c/span\u003ed) confirm the stable and excellent corrosion resistance of both HSP and LSP samples. Notably, the pitting potential remains consistently high at approximately 1000 mV, unaffected by variations in SP. This exceptional pitting resistance, coupled with ultra-high tensile strength, positions our material significantly above all other reference materials, including DSSs, In718, C276 and CoCrFeNiTi high entropy alloys [\u003cspan citationid=\"CR36\" class=\"CitationRef\"\u003e36\u003c/span\u003e, \u003cspan citationid=\"CR37\" class=\"CitationRef\"\u003e37\u003c/span\u003e]. This synergistic combination of properties unequivocally indicates that the HSP steel achieves a breakthrough in simultaneously enhancing both strength and corrosion resistance.\u003c/p\u003e \u003cp\u003eIn summary, this study conclusively demonstrates that interface velocity-driven non-equilibrium nitrogen supersaturation during additive manufacturing provides an efficient and universal strategy to overcome the long-standing strength-corrosion trade-off in metallic alloys. The engineered non-equilibrium microstructure, characterized by ultra-high tensile strength (~\u0026thinsp;1700 MPa) and exceptional pitting corrosion resistance (~\u0026thinsp;1000 mV), decisively breaks the classical strength-corrosion trade-off. This work offers a powerful new pathway for designing hierarchical microstructures and accelerating the rational development of next-generation high-performance alloys through additive manufacturing.\u003c/p\u003e \u003cp\u003e \u003c/p\u003e"},{"header":"Discussion","content":"\u003cdiv id=\"Sec8\" class=\"Section2\"\u003e \u003ch2\u003eContribution of non-equilibrium precipitation for strength increase\u003c/h2\u003e \u003cp\u003e \u003c/p\u003e\u003cdiv class=\"BlockQuote\"\u003e \u003cp\u003eTo quantify strengthening mechanisms, we first establish a baseline for LPBF-processed high nitrogen ferritic stainless steels without extensive precipitation hardening. Nie et al. [\u003cspan class=\"CitationRef\"\u003e38\u003c/span\u003e] reported a ~ 800 MPa yield strength in LPBF-fabricated duplex stainless steel (predominantly ferritic, ~ 96%), attributed to grain refinement, solid solution strengthening, and high geometrically necessary dislocation (GND) density (~ 1.4 × 10\u003csup\u003e14\u003c/sup\u003e m\u003csup\u003e-2\u003c/sup\u003e). He et al. [\u003cspan class=\"CitationRef\"\u003e3\u003c/span\u003e] corroborated this, showing dislocation density saturates at 10\u003csup\u003e13\u003c/sup\u003e-10\u003csup\u003e14\u003c/sup\u003e m\u003csup\u003e-2\u003c/sup\u003e for scan speeds ≥ 700 mm/s, with minimal laser power impact. This consistent high dislocation density (\u003cb\u003eFig.\u0026nbsp;2c\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e) is a fundamental characteristic of rapid solidification LPBF, forming a reliable baseline strength.\u003c/p\u003e \u003cp\u003eAssuming similar high dislocation densities (~ 10\u003csup\u003e14\u003c/sup\u003e m⁻²) for our LSP and HSP conditions, the dislocation strengthening (\u003cb\u003eΔσ\u003c/b\u003e\u003csub\u003e\u003cb\u003edis\u003c/b\u003e\u003c/sub\u003e.) is comparable. Using the Taylor hardening model [\u003cspan class=\"CitationRef\"\u003e39\u003c/span\u003e] with ferritic parameters (Taylor factor M = 2.7, shear modulus G = 83 GPa, Burgers vector b = 0.25 nm, and obstacle efficiency α = 0.475), \u003cb\u003eΔσ\u003c/b\u003e\u003csub\u003e\u003cb\u003edis.\u003c/b\u003e\u003c/sub\u003e is estimated at 315 MPa. For solid solution strengthening (Δσ\u003csub\u003eSS\u003c/sub\u003e), the empirical relationship Δσ\u003csub\u003eSS\u003c/sub\u003e = 1103.45C + 1103.45N + 25.8Si + 19.2Ni + 16.9Mn + 15.9Mo + 2.6Cr (in MPa, with element concentrations in wt. %) [\u003cspan class=\"CitationRef\"\u003e40\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e41\u003c/span\u003e] was applied. Adopting a realistic nitrogen content of 0.1 wt.% (\u003cb\u003eFig. S8a\u003c/b\u003e), due to its low equilibrium solubility in bcc-Fe, Δσ\u003csub\u003eSS\u003c/sub\u003e totals 439 MPa. Grain refinement strengthening (Δσ\u003csub\u003ehp\u003c/sub\u003e) was quantified via the Hall-Petch relationship [\u003cspan class=\"CitationRef\"\u003e42\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e43\u003c/span\u003e]. With a Hall-Petch coefficient (k\u003csub\u003ey\u003c/sub\u003e) of 0.55 MPa·m\u003csup\u003e1/2\u003c/sup\u003e [\u003cspan class=\"CitationRef\"\u003e44\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e45\u003c/span\u003e] and measured average grain sizes of 26.1 µm (LSP) and 6.2 µm (HSP), Δσ\u003csub\u003ehp\u003c/sub\u003e values are approximately 108 MPa and 220 MPa, respectively. \u003cb\u003eTable S3\u003c/b\u003e summarizes the quantitative contributions to yield strength. This highlights that non-equilibrium CrN precipitation, activated at higher scan speed, is the dominant mechanism, accounting for ~ 72% of the total strength increment and critically contributing to superior mechanical properties for HSP sample.\u003c/p\u003e \u003c/div\u003e \u003cp\u003e\u003c/p\u003e \u003c/div\u003e\n\u003ch3\u003eInterface velocity-driven interstitial solute trapping accounting for counterintuitive enhanced precipitation\u003c/h3\u003e\n\u003cp\u003eConventional wisdom dictates that slower cooling rates favor precipitation by allowing sufficient time for solute diffusion and phase formation, a principle well-established in phase transformation kinetics and observed in traditional alloy processing [\u003cspan class=\"CitationRef\"\u003e46\u003c/span\u003e–\u003cspan class=\"CitationRef\"\u003e48\u003c/span\u003e]. Similarly, in additive manufacturing, lower scan speeds (implying slower cooling) often facilitate detrimental precipitate formation [\u003cspan class=\"CitationRef\"\u003e5\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e49\u003c/span\u003e–\u003cspan class=\"CitationRef\"\u003e51\u003c/span\u003e]. However, this study reveals a counterintuitive finding: increasing laser scan speed significantly enhances nanoscale CrN precipitation, as evidenced by the remarkably higher density of nanoscale precipitates in the HSP sample compared to the LSP (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003e) and its contribution to strength (see above subsection). This \"higher scan speed, more CrN precipitation\" phenomenon (illustrated in \u003cb\u003eFig. S9\u003c/b\u003e) reveals substantial nitrogen enrichment in the ferritic matrix, significant grain refinement, dramatically increased density of intragranular CrN nanoprecipitates, and suppressed PFZ under high scan speed, while maintaining a continuous corrosion-resistant grain boundary austenite network.\u003c/p\u003e \u003cp\u003eThis anomalous precipitation behavior is driven by the unique non-equilibrium solidification and thermal cycles inherent to LPBF (Figs.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003e, S8). The fundamental mechanism involves two key aspects. Firstly, accelerated solidification, primarily controlled by the solid-liquid interface velocity and regulated by laser scan speed, leads to \"solid-liquid interface velocity dominating nitrogen supersaturation\" as supported by our simulations (Figs.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ec, \u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ed). Secondly, accelerated solidification dramatically suppresses long-range diffusion of interstitial nitrogen, kinetically trapping a vastly higher concentration of solutes within the ferritic matrix, forming a highly supersaturated solid solution. The resultant enormous chemical driving force (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003eb) then promotes the prolific, homogeneous nucleation of fine, coherent metastable CrN precipitates, favored over Cr\u003csub\u003e2\u003c/sub\u003eN due to its lower coherent interfacial energy with the ferrite matrix (\u003cb\u003eFig.\u0026nbsp;2d\u003c/b\u003e\u003csub\u003e\u003cb\u003e1\u003c/b\u003e\u003c/sub\u003e). These fine precipitates contribute substantially to the ~ 600 MPa strength increase \u003cb\u003e(Table S3\u003c/b\u003e) in the HSP sample. This is further facilitated by the rapid thermal cycles of LPBF, which generate substantial thermal stress and plastic strain, leading to a high density of dislocations (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e2\u003c/span\u003ec) and vacancies [\u003cspan class=\"CitationRef\"\u003e52\u003c/span\u003e–\u003cspan class=\"CitationRef\"\u003e54\u003c/span\u003e] that act as preferential nucleation sites, thereby enhancing precipitation nucleation.\u003c/p\u003e \u003cp\u003eTherefore, this counterintuitive enhanced precipitation is a synergistic outcome of interface velocity-driven ultra-fast solidification, which simultaneously creates (i) high supersaturation (thermodynamic driving force) and (ii) high defect density (kinetic pathway), collectively facilitating a burst of fine-scale precipitation otherwise kinetically inaccessible.\u003c/p\u003e\n\u003ch3\u003eUltra-fast solidification induced nitrogen-deficient austenite and associated toughening via twinning\u003c/h3\u003e\n\u003cp\u003eCommercial 2507/3207 high-nitrogen duplex stainless steels exhibit extreme austenite stability, precluding deformation-induced twinning or martensitic transformation [\u003cspan class=\"CitationRef\"\u003e48\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e55\u003c/span\u003e]. Our LPBF process overcomes this by engineering a metastable grain boundary austenite, activating additional energy-dissipating deformation mechanisms. LPBF's inherent non-equilibrium solidification is key. Extreme solidification rates kinetically trap nitrogen in the ferritic matrix, suppressing its partitioning to austenite. Consequently, grain boundary austenite is significantly nitrogen-depleted (~ 1 wt.% N, Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e4\u003c/span\u003ea) compared to conventional counterparts. This compositional shift is critical, as nitrogen potently modulates austenite stacking fault energy (SFE) [\u003cspan class=\"CitationRef\"\u003e56\u003c/span\u003e]. While not monotonic, DFT calculations [\u003cspan class=\"CitationRef\"\u003e57\u003c/span\u003e] suggest that increasing nitrogen to 1.0 wt.% promotes twinning over cross-slip, consistent with our observations (Fig.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003ee). This nitrogen-deficient state substantially lowers SFE, rendering the boundary austenite mechanically metastable.\u003c/p\u003e \u003cp\u003eBeyond deformation twins, a long-range ordered 9R structure was observed at twin boundaries (\u003cb\u003eFig.\u0026nbsp;3e\u003c/b\u003e\u003csub\u003e\u003cb\u003e3\u003c/b\u003e\u003c/sub\u003e). Both nanoscale twins and 9R structures act as potent obstacles to dislocation motion, enhancing work hardening (TWIP effect) [\u003cspan class=\"CitationRef\"\u003e29\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e58\u003c/span\u003e]. Simultaneously, grain boundary austenite impedes crack propagation via ductile bridging and blunting [\u003cspan class=\"CitationRef\"\u003e59\u003c/span\u003e]. Thus, toughening arises from a process-enabled synergy: LPBF-induced N-depletion creates low-SFE metastable austenite, activating stacking faulting and deformation twinning. This, combined with austenite's metastable characteristics and enhanced CrN precipitation, comprehensively explains the excellent strength-ductility synergy under HSP condition. Our work highlights that in AM processes under extreme non-equilibrium, traditional precipitation kinetics can be inverted, enabling novel microstructural design via precise thermal kinetic control.\u003c/p\u003e \u003cdiv id=\"Sec11\" class=\"Section2\"\u003e \u003ch2\u003eNon-equilibrium solidification engineering grain boundary nitrogen for boosted corrosion resistance​\u003c/h2\u003e \u003cp\u003eThis study reveals exceptional pitting corrosion resistance (Epit ≈ 1000 mV) HSP sample, despite a high density of chromium-rich nitride precipitates. This contradicts the established paradigm where CrN precipitation typically degrades localized corrosion resistance due to adjacent chromium depletion [\u003cspan class=\"CitationRef\"\u003e60\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e61\u003c/span\u003e]. The anomaly stems from a unique, multi-scale corrosion protection architecture engineered by LPBF.\u003c/p\u003e \u003cp\u003eFirst, at grain boundaries (GBs), the refined microstructure avoids vulnerable corrosion pathways. TEM analysis shows the GBs are either enriched with nitrogen-stabilized austenite or decorated with nanoscale, continuous Cr, N-rich layers (Figs.\u0026nbsp;\u003cspan class=\"InternalRef\"\u003e3\u003c/span\u003e, S5 and S6). This configuration, particularly Cr, N enrichment in austenite, results in a localized pitting resistance equivalent (PREN = %Cr + 3.3 × %Mo + 16 × %N) [\u003cspan class=\"CitationRef\"\u003e62\u003c/span\u003e] significantly higher than the ferritic grain interior. Consequently, these GBs act as corrosion-resistant barriers, suppressing intergranular attack and hindering pit propagation.\u003c/p\u003e \u003cp\u003eSecond, within ferritic grains, the high-density nanoprecipitates differ fundamentally from coarse, equilibrium nitrides. Formed under extreme non-equilibrium solidification, these precipitates are ultra-fine and exhibit coherent/semi-coherent interfaces (\u003cb\u003eFig.\u0026nbsp;2d\u003c/b\u003e\u003csub\u003e\u003cb\u003e2\u003c/b\u003e\u003c/sub\u003e). Crucially, the surrounding chromium-depleted zone is extremely narrow and chemically mild due to suppressed diffusion and non-equilibrium composition (\u003cb\u003eFig. S8c\u003c/b\u003e). The adjacent matrix's chromium concentration remains sufficiently high to maintain passivity [\u003cspan class=\"CitationRef\"\u003e7\u003c/span\u003e, \u003cspan class=\"CitationRef\"\u003e63\u003c/span\u003e]. Thus, these intragranular nanoprecipitates, while providing potent Orowan strengthening, do not create the extensive, interconnected chromium-depleted networks that initiate stable pitting, unlike coarse counterparts in conventionally processed materials.\u003c/p\u003e \u003cp\u003eIn summary, high corrosion resistance is preserved through a synergistic, multi-level defense: (1) corrosion-resistant GBs with high localized PREN, and (2) intragranular nanoprecipitates with minimal detrimental impact on the local electrochemical environment. This unique microstructural configuration, enabled by rapid solidification and defect-assisted precipitation kinetics of high-speed LPBF, successfully decouples nitrides' strengthening effect from their traditionally deleterious impact on corrosion resistance.\u003c/p\u003e \u003c/div\u003e \u003cdiv id=\"Sec12\" class=\"Section2\"\u003e \u003cdiv id=\"Sec13\" class=\"Section3\"\u003e \u003c/div\u003e \u003c/div\u003e "},{"header":"Methods","content":"\u003ch2\u003ePowder preparation \u0026amp; LPBF processing\u003c/h2\u003e\n\u003cp\u003eThe high nitrogen stainless steel powder for LPBF printing was prepared by vacuum induction melting gas atomization (VIGA) method, which was finished in the Institute of Metal Research, Chinese Academy of Sciences. The chemical compositions of the powder were determined by various methods (ICP method, ONH and CS analyzers) as shown in \u003cstrong\u003eTable S1\u003c/strong\u003e. This composition is characterized by a high nitrogen content, approximately 0.6 wt.%, hence the LPBF steel is conveniently named 6N. After powder preparation, the powder undergoes further sieving to select powder with a particle size of approximately 15\u0026ndash;53 \u0026micro;m for printing. The powders were dried for at least 5 h in a vacuum oven at 70 ℃ before additive manufacturing fabrication.\u003c/p\u003e\n\u003cp\u003eThe LPBF processing was performed via a 3D additive manufacturing system (Hans-M-100, Hans Laser) equipped with a continuous Ytterbium fiber laser (wavelength: 1064 nm, spot diameter: 50 \u0026micro;m). Numerous orthogonal experiments (hatch spacing: 80\u0026ndash;120 \u0026micro;m, scanning speed: 200\u0026ndash;1500 mm/s, laser power 200\u0026ndash;300 W, layer thickness: 30\u0026ndash;60 \u0026micro;m) have been explored previously, and a relatively optimal printing parameter space with varying scanning speed has been obtained, as presented in \u003cstrong\u003eTable S2\u003c/strong\u003e. Apart from the scanning speed, the printing power (\u003cstrong\u003eP\u003c/strong\u003e), layer thickness (\u003cstrong\u003et\u003c/strong\u003e), hatch spacing (\u003cstrong\u003eH\u003c/strong\u003e), and printing strategy remain unchanged at \u003cstrong\u003eP\u003c/strong\u003e\u0026thinsp;=\u0026thinsp;250 W, \u003cstrong\u003et\u003c/strong\u003e\u0026thinsp;=\u0026thinsp;50 \u0026micro;m, \u003cstrong\u003eH\u003c/strong\u003e\u0026thinsp;=\u0026thinsp;80 \u0026micro;m, respectively. The commonly bidirectional scanning strategy with a rotation angle of 67\u003csup\u003eo\u003c/sup\u003e between the upper and lower layers was used for minimizing texture and anisotropy as shown in \u003cstrong\u003eFig. S1a\u003c/strong\u003e. All samples were fabricated in a high-purity flowing argon atmosphere for avoiding oxidation or nitriding. A 316L-type stainless steel substrate (110 mm \u0026times; 110 mm \u0026times; 13 mm) was used due to its compatibility with stainless steel materials.\u003c/p\u003e\n\u003ch2\u003ePlate tensile \u0026amp; potentiodynamic polarization tests\u003c/h2\u003e\n\u003cp\u003eThese cuboid samples with different scanning parameters were cut from the 316L substrate and were further cut into several plate tensile samples using wire electrical discharge machining. The plane of the tensile sample was parallel to the substrate and other sampling details like sample size and thickness are shown in \u003cstrong\u003eFig. S1b\u003c/strong\u003e. All tensile samples were mechanically ground using SiC sandpapers ranging from 180 to 2000 grit to obtain flat and clean surfaces. Tensile tests were conducted on a universal testing machine equipped with an optical extensometer. The tensile strain rate was 10\u003csup\u003e\u0026minus;\u0026thinsp;3\u003c/sup\u003e s\u003csup\u003e\u0026minus;\u0026thinsp;1\u003c/sup\u003e.\u003c/p\u003e\n\u003cp\u003eAfter tensile testing, the clamping planes of the fractured samples were used for the potentiodynamic polarization experiment. All samples were mechanically re-polished using silicon carbide paper with grit sizes ranging from 180 to 2000 and finished by using 2.5 \u0026micro;m diamond suspension for 30 minutes and vibratory polishing for 2 hours. The electrochemical measurements were conducted on an electrochemical workstation (Versa STAT 3 F) using a three-electrode cell. The working electrode was an investigated sample that was fixed in a holder with an exposing area of ~\u0026thinsp;0.28 cm\u003csup\u003e2\u003c/sup\u003e, a platinum plate was used as the counter electrode, and a saturated Ag/AgCl electrode was used as the reference electrode.\u003c/p\u003e\n\u003cp\u003ePotentiodynamic polarization tests were carried out in 3.5% NaCl electrolyte, with a standard scan rate of 0.167 mV/s, an initial potential of \u0026minus;\u0026thinsp;0.6 V, and a final potential of 1.2 V against the open circuit potential (OCP). The potentiodynamic polarization test and electrochemical resistance spectroscopy (EIS) were conducted after OCP by immersing the sample for 30 min in the electrolyte to wait for the system to stabilize. EIS measurements were performed by using a frequency range from 10\u003csup\u003e5\u003c/sup\u003e Hz to 10\u003csup\u003e\u0026minus;\u0026thinsp;2\u003c/sup\u003e Hz, and the root mean square (RMS) potential amplitude is 10 mV. The electrochemical measurements were conducted at room temperature (22\u0026ndash;24\u0026deg;C).\u003c/p\u003e\n\u003ch2\u003eMicrostructural characterization\u003c/h2\u003e\n\u003cp\u003eThe morphologies and crystal orientation of these 6N samples were observed by a field emission scanning electron microscope (SEM, Apreo 2, FEI) equipped with a high-speed electron backscatter diffraction (EBSD, Velocity, EDAX) detector and a scanning transmission electron microscope (STEM) detector. EBSD maps were obtained with varying step sizes of 0.05 and 0.1 \u0026micro;m according to collection region areas. Raw EBSD data were analyzed using the OIM (EDAX) software.\u003c/p\u003e\n\u003cp\u003eFor TEM sampling, both FIB and electrochemical jet polishing methods were used. Some 3 mm disks were twin-jet electro-polished by a TenuPol-5 (Struers) in an 8% HClO\u003csub\u003e4\u003c/sub\u003e ethanol solution at 25\u0026deg;C and 20 V. The TEM lamellas were prepared via FIB (Helios 600i, FEI) with Pt protection, coarse milling at 30 kV, and final polishing at 5 kV to ensure electron transparency. A 200 kV transmission electron microscope (TEM, Talos F200X G2, FEI) equipped with high-angle annular dark-field scanning transmission electron microscopy (HAADF-STEM) functionality, and energy-dispersive X-ray spectrometer (EDS) was employed for microstructural analysis.\u003c/p\u003e\n\u003ch2\u003eSimulation methods\u003c/h2\u003e\n\u003cp\u003eA three-dimensional thermo-fluidic simulation of the single-track melt-pool was performed. The model solves the governing equations of mass, momentum, and energy conservation to predict the transient temperature field, melt-pool geometry, and thermal history under different laser scan speeds. A Gaussian volumetric heat source was employed to simulate the laser energy input. Temperature-dependent thermophysical properties (e.g., thermal conductivity, specific heat, viscosity) of the alloy were incorporated from the Thermo-Calc FEDEMO V5 database. The simulations were conducted for different scanning speed, corresponding to the experimental conditions.\u003c/p\u003e\n\u003cp\u003eThermodynamic and kinetic calculations were employed to understand the phase stability and solute redistribution behavior under the extreme non-equilibrium solidification conditions of LPBF. Equilibrium phase diagram calculations and driving force analyses were performed using the Thermo-Calc software with a relevant thermodynamic database (TCFE7). A vertical section of the multi-component system (Fe-Cr-Ni-Mo-Si-N) was calculated to map the stability regions of ferrite (BCC) and austenite (FCC) as a function of nitrogen content and temperature. Furthermore, the normalized driving forces for the precipitation of Cr\u003csub\u003e2\u003c/sub\u003eN and CrN phases from the ferrite matrix were calculated based on its nominal composition. Diffusion simulations were performed using the Dictra module. A one-dimensional diffusion couple model was constructed, consisting of ferrite (BCC_A2), Cr\u003csub\u003e2\u003c/sub\u003eN (HCP) and austenite (FCC_A1) phases. Simulations modeled the continuous cooling from 1360\u003csup\u003eo\u003c/sup\u003eC at rates representative of the melt pool simulation results.\u003c/p\u003e"},{"header":"Declarations","content":"\u003cp\u003e\u003cstrong\u003eData availability\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eData will be made available on request.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eDeclaration of competing interest\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eThe authors declare that they have no known competing financial interests or personal relationships that could have appeared to influence the work reported in this paper.\u003c/p\u003e\n\u003cp\u003e\u003cstrong\u003eAcknowledgements\u003c/strong\u003e\u003c/p\u003e\n\u003cp\u003eM. Yang acknowledges the support from the Guangdong Basic and Applied Basic Research Foundation (Grant No. 2023A1515240048), the Shenzhen Science and Technology Program (Grant No. JCYJ20240813094806009) and the National Natural Science Foundation of China (Grant No. 52301153); J. Yi acknowledges the support from the National Natural Science Foundation of China (Grant No. 52301154) and the Shenzhen Science and Technology Program (Grant No. JCYJ20220530114400001); M. He acknowledges the support from the Guangdong Basic and Applied Basic Research Foundation (Grant No. 2023A1515110595); Z. Yao acknowledges the support from the National Natural Science Foundation of China (Grant No. 52301042) and the Natural Science Foundation of Top Talent of SZTU (grant No. GDRC202532). All authors gratefully acknowledge the assistance from Dr. Y. Qiu and Dr. D. He at SUStech Core Research Facilities and the electron microscope center of KAIPLE Co. Ltd (Changsha) for the support of microstructural characterizations.\u003c/p\u003e"},{"header":"References","content":"\u003col\u003e\n \u003cli\u003eS.F. Nabavi, H. Dalir, A. Farshidianfar, A comprehensive review of recent advances in laser powder bed fusion characteristics modeling: metallurgical and defects, The International Journal of Advanced Manufacturing Technology 132(5) (2024) 2233\u0026ndash;2269.\u003c/li\u003e\n \u003cli\u003eN. Ren, J. Li, R. Zhang, C. Panwisawas, M. Xia, H. Dong, J. Li, Solute trapping and non-equilibrium microstructure during rapid solidification of additive manufacturing, Nature communications 14(1) (2023) 7990.\u003c/li\u003e\n \u003cli\u003eX. He, V. Rielli, Q. Liu, X. Li, V. Luzin, N. Haghdadi, S. Primig, Effects of laser powder bed fusion parameters on the delta-ferrite to austenite phase transformation in duplex stainless steels, Additive Manufacturing (2025) 104825.\u003c/li\u003e\n \u003cli\u003eY. Fang, M.-K. Kim, Y. Zhang, T. Kim, J. No, J. 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Pan, Study of corrosion behavior of a 2507 super duplex stainless steel: influence of quenched-in and isothermal nitrides, International Journal of Electrochemical Science 9(1) (2014) 61\u0026ndash;80.\u003c/li\u003e\n\u003c/ol\u003e"}],"fulltextSource":"","fullText":"","funders":[],"hasAdminPriorityOnWorkflow":false,"hasManuscriptDocX":true,"hasOptedInToPreprint":true,"hasPassedJournalQc":"","hasAnyPriority":true,"hideJournal":true,"highlight":"","institution":"","isAcceptedByJournal":false,"isAuthorSuppliedPdf":false,"isDeskRejected":"","isHiddenFromSearch":false,"isInQc":false,"isInWorkflow":false,"isPdf":false,"isPdfUpToDate":true,"isWithdrawnOrRetracted":false,"journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true},"keywords":"","lastPublishedDoi":"10.21203/rs.3.rs-9249060/v1","lastPublishedDoiUrl":"https://doi.org/10.21203/rs.3.rs-9249060/v1","license":{"name":"CC BY 4.0","url":"https://creativecommons.org/licenses/by/4.0/"},"manuscriptAbstract":"\u003cp\u003eAchieving synergistic high strength and corrosion resistance in high-nitrogen stainless steels remains challenging due to inherent trade-offs: dispersion strengthening enhances strength, but heterogeneous interfaces often act as micro-galvanic corrosion or pitting nucleation sites, compromising corrosion resistance. Laser powder bed fusion (LPBF) with its extreme thermal cycles, offers a non-equilibrium strategy to address this limitation. Here, using multi-scale microstructural characterization and thermo-kinetic simulations methods, we report that during melt pool solidification, interface velocity-regulated kinetics trap excess nitrogen in the ferritic matrix, inducing non-equilibrium nitrogen supersaturation. Subsequently during fast cooling process, this supersaturated nitrogen combines with chromium to form dense superfine CrN lamella. Coupled with fine grain structures from rapid solidification, this two-stage non-equilibrium process achieves a synergy of high tensile strength (~1700 MPa) and pitting resistance (~1000 mV), breaking the classical strength-corrosion trade-off. Key mechanisms include: (1) superfine CrN lamellae minimizing adjacent Cr-depleted zones, (2) grain refinement suppressing precipitate-free zones (PFZ), and (3) nitrogen-stabilized grain boundaries enhancing corrosion resistance.\u003c/p\u003e\n\u003cp\u003eBy controlling interface velocity in LPBF non-equilibrium solidification, we establish a nitrogen supersaturation pathway that tailors hierarchical microstructures (grains, boundaries, nanoprecipitates), resolving the strength-corrosion trade-off in additively manufactured high-nitrogen alloys.\u003c/p\u003e","manuscriptTitle":"Interface velocity-driven non-equilibrium nitrogen supersaturation in additive manufacture: a universal strategy for breaking strength-corrosion trade off","msid":"","msnumber":"","nonDraftVersions":[{"code":1,"date":"2026-04-09 02:20:58","doi":"10.21203/rs.3.rs-9249060/v1","editorialEvents":[{"type":"communityComments","content":0}],"status":"published","journal":{"display":true,"email":"[email protected]","identity":"researchsquare","isNatureJournal":false,"hasQc":true,"allowDirectSubmit":true,"externalIdentity":"","sideBox":"","snPcode":"","submissionUrl":"/submission","title":"Research Square","twitterHandle":"researchsquare","acdcEnabled":true,"dfaEnabled":false,"editorialSystem":"","reportingPortfolio":"","inReviewEnabled":false,"inReviewRevisionsEnabled":true}}],"origin":"","ownerIdentity":"55acc85f-bc17-4c1d-97c2-50b834562110","owner":[],"postedDate":"April 9th, 2026","published":true,"recentEditorialEvents":[],"rejectedJournal":[],"revision":"","amendment":"","status":"posted","subjectAreas":[{"id":65420959,"name":"Physical sciences/Materials science/Structural materials/Metals and alloys"},{"id":65420960,"name":"Physical sciences/Materials science/Structural materials/Mechanical properties"}],"tags":[],"updatedAt":"2026-04-09T20:00:16+00:00","versionOfRecord":[],"versionCreatedAt":"2026-04-09 02:20:58","video":"","vorDoi":"","vorDoiUrl":"","workflowStages":[]},"version":"v1","identity":"rs-9249060","journalConfig":"researchsquare"},"__N_SSP":true},"page":"/article/[identity]/[[...version]]","query":{"redirect":"/article/rs-9249060","identity":"rs-9249060","version":["v1"]},"buildId":"XKTyCvWXoU3ODBz1xrDgd","isFallback":false,"isExperimentalCompile":false,"dynamicIds":[84888],"gssp":true,"scriptLoader":[]}