Achieving Homogeneous Microstructure and Superior Properties in High-N Austenitic Stainless Steel via a Novel Atmosphere-Switching Method

Abstract

Powder metallurgy is widely used to fabricate high-nitrogen, nickel-free austenitic stainless steel. However, after sintering and nitriding, additional solution treatment is typically required to achieve uniform nitrogen distribution and a homogeneous austenite phase. This work proposes a novel method to eliminate the need for lengthy and high-temperature solution treatment by switching the nitrogen atmosphere to argon during the cooling process. The effects of different N₂-Ar atmosphere-switching temperatures (750–1320 °C) on the phase composition, element distribution, microstructure, mechanical properties, and corrosion resistance of the studied steels were systematically investigated. Results show that cooling in the N₂ atmosphere initially transforms the matrix to a fully austenitic structure enriched with nitrogen. Excessive nitrogen infiltration leads to Cr₂N precipitation, inducing partial austenite decomposition and forming a multiphase structure comprising austenite, α-Fe, and Cr₂N. Strategic switching from N₂ to Ar reverses this reaction, yielding a high-nitrogen, chemically uniform austenitic structure. Specifically, switching at 1150 °C, the steel exhibits excellent mechanical properties and corrosion resistance, with a yield strength of 749 MPa, an ultimate tensile strength of 1030 MPa, an elongation of 38.7%, and a corrosion current of 0.06 mA/cm², outperforming the steels cooled solely in N₂ and subsequently solution-treated. This novel method offers advantages in cost reduction, energy saving, and operational effectiveness, highlighting its potential for broad industrial application.

1. Introduction

Nitrogen (N), as an alloying element, is commonly integrated into stainless steel due to its capacity to significantly enhance both mechanical properties and corrosion resistance [1,2]. Currently, two primary methods for incorporating nitrogen into stainless steel are gas–liquid nitriding and gas–solid nitriding [3,4]. Gas–liquid nitriding involves infusing nitrogen into molten steel. However, the solubility of nitrogen in liquid steel is notably low, at only 0.045 wt.%. Achieving high nitrogen content in such an environment requires operating at high pressures, resulting in increased equipment costs and safety risks. Conversely, gas–solid nitriding presents a safer and more reliable alternative; however, the nitriding rate is sluggish, particularly for large-scale steel components, posing challenges for nitrogen permeation in the steel interior [5,6,7,8].

Powder metallurgy technology has garnered significant attention as a potential solution to this challenge [9,10,11,12,13,14]. The high porosity of raw billets facilitates nitrogen diffusion. During heating and sintering, nitrogen can permeate raw billets, effectively leading to high nitrogen content throughout. However, the direct contact between the steel surface and N₂ accelerates nitriding on the surface, resulting in higher nitrogen concentrations in the surface layer. Consequently, there persists a variation in nitrogen concentration along the cross-section of the steel parts. This discrepancy is less noticeable before sintering at high temperatures due to the low density of raw billets (generally 60–70%), allowing nitrogen easy access. During sintering (above 1200 °C), the discrepancy remains minor as nitrogen diffusion rates increase substantially [15,16,17,18]. However, the disparity becomes markedly pronounced during cooling, as the density of steel billets rises (generally above 90%) and nitrogen diffusion decelerates. Consequently, nitriding predominantly occurs on the surface layer during cooling. This variance frequently yields differences in microstructures and properties across the steel cross-section, thereby limiting its performance and applicability.

The conventional approach to address this issue involves solution treatment [19,20,21,22,23,24,25], which requires reheating the steels to over 1000 °C, holding for a specific duration, and then quenching. In this paper, a novel method is proposed to achieve uniform chemical composition, a microstructure predominantly comprising a single austenite phase, and desirable properties without resorting to solution treatment. Instead, the novel method involves altering the atmosphere at an appropriate temperature during cooling to facilitate the escape of excess nitrogen from the surface layer.

2. Materials and Methods

The steels were prepared using metal injection molding (MIM) process. N₂ gas atomized duplex nickel-free stainless steel powder (provided by Sandvik Osprey Ltd., Neath, UK) was used as raw materials. The chemical composition of the powder is detailed in

Table 1. Chemical composition of the powder.

Chemical Composition	Cr	Mn	Mo	N	C	Si	P	S	Ni	Fe
wt.%	16.9	11	3.28	0.4	0.07	0.7	0.013	0.005	0.1	Balance

. The high content of manganese serves to facilitate the solid solution of nitrogen and enhance the stability of austenite, thereby offering a cost-effective alternative to nickel in conventional austenitic stainless steel. Additionally,

Table 2. Composition and properties of binder systems.

Composition	Melting Point/°C	Decomposition Temperature Range/°C	Density/g∙cm−3	Mass Fraction in Binder/wt.%
Microcrystalline wax	75	300–500	0.920	65
High-density polyethylene	133.5	420–510	0.964	25
Ethylene–vinyl acetate	86	310–500	0.941	5
stearic acid	67	290–370	0.941	5

presents the composition and properties of the binder utilized in our experiment, with the volume fraction of binder in the feed set at 32%.

The preparation process involved mixing, injection, debinding, and sintering. Mixing was conducted at 140 °C with a speed of 30 revolutions per minute (r/min). Subsequently, injection was performed at 160 °C under a pressure of 100 bar. Following this step, cylindrical samples with a diameter of 25 mm and a thickness of 6 mm were obtained. For debinding, a two-step procedure was employed. Initially, microcrystalline wax was dissolved in n-heptane at 62 °C for 6 h, followed by the removal of the remaining binder through thermal debinding under flowing N₂. Sintering and solution treatment were carried out in a tube furnace, with the heating and cooling curves depicted in Figure 1. Based on our previous research results, a relatively high sintering temperature of 1320 °C was chosen in Figure 1a to achieve higher density and strength without inducing liquid phase sintering. In Figure 1b, a solution treatment temperature of 1150 °C was selected to balance solubility efficiency and effectiveness. Detailed sample labels and their corresponding process parameters are provided in

Table 3. Sample codes and corresponding processing parameters.

Sample Label	Key Processing Steps
750Ar	S-N process → Switch to Ar atmosphere when cooled to 750 °C
950Ar	S-N process → Switch to Ar atmosphere when cooled to 950 °C
1150Ar	S-N process → Switch to Ar atmosphere when cooled to 1150 °C
1320Ar	S-N process → Switch to Ar atmosphere directly at 1320 °C
1320N2	S-N process → Keep N2 atmosphere throughout during cooling
1320N2-ST	As-prepared 1320N2 steel → Solution-treated in N2 atmosphere

Note: “S-N process” means sintering and nitriding at 1320 °C for 2 h.

. In these labels, the numeric value at the beginning indicates the Celsius temperature when the atmosphere switches, with “N2” and “Ar” representing the atmosphere, and “ST” denoting solid solution treatment. Of note, this work proposes a novel process termed the “Sintering-Nitriding-Atmosphere Transition (SNAT)” process, wherein, after sintering and nitriding of the prefabricated billets, an atmospheric transition occurs during the cooling process. The duration of the atmosphere-switching process lasts approximately 1 min, after which the flow of the atmosphere is sustained. This method enables the direct attainment of a uniform microstructure with high-nitrogen content and a single-phase austenitic structure while eliminating the need for solid solution treatment in the conventional process.

Microstructure investigation was conducted using both an optical microscope (OM) and a scanning electron microscope (SEM, Sigma300, ZEISS, Oberkochen, Germany), and the acceleration voltage is 15 kV using the secondary electron probe mode. Before the examination, the polished surface of the samples underwent chemical etching in a solution comprising 75 vol.% HCl and 25 vol.% HNO₃ for approximately 15 s. SEM analysis, equipped with an energy-dispersive X-ray spectroscope (EDS, Ultim Extreme, Oxford, UK), was employed to analyze the chemical composition. Phase identification was carried out using X-ray diffraction (XRD, X’pert, Panalytical, Almelo, The Netherlands) operating at 40 kV and 40 mA with Cu-Kα radiation, and XRD patterns were recorded with a scanning speed of 12°/min.

Tensile properties were measured using a universal testing machine (MTS Test Star 810, MTS Systems Corporation, Eden Prairie, USA). Tensile test specimens, with a gauge length of 8 mm, width of 2 mm, and thickness of 1 mm, were machined from as-prepared bulk samples after removing the surface layer. Tensile tests were performed at a constant rate of 0.0083 mm/s, corresponding to an initial strain rate of 1 × 10⁻³ s⁻¹. Vickers hardness was measured using a force of 0.98 N (HV0.1) with a dwell time of 15 s, and the hardness was calculated by averaging 15 measurements.

The electrochemical performance was tested by an electrochemical workstation (IviumStat, Ivium Technologies BV, Eindhoven, The Netherlands) at a temperature of 25 ± 0.5 °C. The test electrolyte consisted of a 3.5 wt.% NaCl solution. A classic three-electrode cell configuration was utilized, with a silver chloride electrode (Ag/AgCl) serving as the reference electrode, a platinum electrode (Pt) as the counter electrode, and the specimen as the working electrode. Before potentiodynamic polarization testing, a stable open circuit potential was attained. During electrochemical characterization, the three electrodes were individually immersed in the electrolytic solution. Polarization curves were generated with a scan rate of 1 mV/s.

3. Results and Discussion

3.1. Effect of SNAT Processing on Phase Structure Evolution

During the high-temperature sintering process, the original composition of α-Fe and austenite gradually transformed into high-temperature δ-Fe and austenite. Figure 2 presents the XRD peak patterns of samples under different process parameters, categorized into two groups for clarity: Group A compares the phase structures under different atmosphere-switching temperatures with the new SNAT process, and Group B compares the phase structures of samples under different process paths.

From Figure 2a, it can be observed that 1320Ar steel displays an α-Fe phase, 1150Ar exhibits a pure austenite phase, 950Ar shows almost complete austenite while 750Ar contains a mixture of austenite, α-Fe, and a small amount of Cr₂N precipitates. During the cooling process, δ-Fe transformed into austenite, and austenite continuously absorbed nitrogen in a nitrogen atmosphere, ultimately forming a single N-enriched austenite. Continuous nitrogen infiltration led to nitrogen supersaturation in the austenite, resulting in the formation of nitrides. Once switched to an Ar atmosphere, nitrogen steadily escaped from the austenite, leading to the reduction of nitrides and subsequent decomposition of the austenite. The nitrogen content in the steel fluctuated during the entire cooling process, akin to ocean tides. After sintering, the 1320Ar steel immediately switched from a nitrogen atmosphere to an Ar atmosphere, causing continuous nitrogen loss throughout the cooling process. The nitrogen deficiency causes the decomposition of austenite, ultimately resulting in the formation of full ferrite. In the case of the 1150Ar steel, the temperature for switching to an Ar atmosphere dropped to 1150 °C. During the cooling process from 1320 °C to 1150 °C, the alloy transformed into full austenite and underwent sufficient nitrogen infiltration, enhancing the stability of N-enriched austenite. Nitride formation could not be ruled out in the sintering neck or grain boundary regions, but during the subsequent switch to the Ar atmosphere for cooling, excessive nitrogen infiltration gradually escapes, leading to nitride reduction and ultimately achieving full austenite at room temperature. Further lowering the temperature for switching to an Ar atmosphere implied a longer time for thorough nitrogen infiltration, with a shorter time for nitrogen escape. Nitride formation increased with decreasing atmosphere-switching temperature, with distinct Cr₂N peaks already appearing in the XRD peak patterns of the 750Ar steel.

Figure 2b shows the XRD results of the 1320N2, 1320N2-ST, and 1150Ar steels. The phase constituents of 1320N2-ST and 1150Ar steels are single austenite phases. The austenitic peaks of 1320N2-ST and 1150Ar steels shift to the left, indicating that the inter-planar spacing of austenite grains is larger than that of 1320N2 steel. The presence of nitrogen in austenite as interstitial atoms will enlarge the crystal plane spacing [26], which indicates that the nitrogen contents in austenite of 1320N2-ST and 1150Ar are higher than that of 1320N2 steel. Further, 1320N2-ST obtains a high concentration of interstitial N atoms by quenching the supersaturated austenite. However, in the absence of rapid cooling, the interstitial N content in 1320N2-Ar austenite is also high. This appears to be due to the Ar atmosphere during cooling, which suppresses the precipitation of nitrides, allowing for more interstitial N retention in austenite.

Figure 3 illustrates the SEM microstructure of steels under various process parameters. As shown in Figure 3a, distinct short-rod precipitates are evident at the grain boundaries of 750Ar steel, corresponding to the Cr₂N phase confirmed by XRD analysis. Besides the inherent susceptibility of grain boundaries to etching, the more crucial reason is the presence of the α and γ phase regions at these boundaries. For 950Ar steel, boundary corrosion is minimal, and the presence of nitrides is significantly reduced, although some remnants persist, as depicted in Figure 3b. As shown in Figure 3c for Sample 1150Ar, the grain boundary interface appears clean, devoid of observed precipitates, with its phase composition entirely austenitic. Precipitates are also absent at the grain boundaries of Sample 1320Ar, as shown in Figure 3d; however, corrosion traces are evident due to its matrix phase being ferritic, which offers less corrosion resistance compared to austenite.

In Figure 3e, consistent with the XRD analysis results, the microstructure of 1320N2 includes austenite, α-Fe, and Cr₂N phases, primarily distributed at the grain boundaries. Its phase composition and distribution microstructure are similar to those of 750Ar steel, albeit with a higher proportion of α-Fe and Cr₂N, indicating more severe austenite decomposition during the cooling process. It can be inferred that austenite at the boundaries likely underwent two consecutive decompositions during cooling. Firstly, the continuous nitrogen diffusion led to nitrogen content surpassing the solid solubility limit of Cr₂N, resulting in the decomposition of supersaturated austenite into austenite and Cr₂N: γ supersaturation → γ and Cr₂N. After the first decomposition, both chromium and nitrogen content in the austenite decreased, further weakening its stability, leading to a second decomposition into ferrite and Cr₂N: γ → α and Cr₂N. Moreover, the solubility of nitrogen in austenite gradually decreased during the cooling process, and since Sample 1320N2 maintained a nitrogen atmosphere throughout the entire cooling process, nitrogen escape was inhibited, thus intensifying the aforementioned decomposition. Given this, nitrided steels often require solution treatment. After solution treatment at 1150 °C for 1.5 h, the Cr₂N dissolved, and the α-Fe transformed back into austenite. As observed in Figure 3f, the microstructure in the interior of 1320N2-ST steel is a single austenite phase.

Interestingly, 1150Ar steel does not require solution treatment but achieves a single austenite structure simply by switching the atmosphere. The conversion of the cooling atmosphere to Ar at 1150 °C causes nitrogen to escape first from supersaturated austenite at boundary regions during cooling. This balances the nitrogen content in austenite, inhibiting the initial decomposition of austenite (γ supersaturated → γ and Cr₂N). Subsequently, because the nitrogen content in austenite is sufficient, the second-step decomposition of austenite (γ → α and Cr₂N) cannot occur, preserving the austenite at room temperature. From this point of view, higher nitrogen content in austenitic steel does not necessarily yield better results [6,27,28,29].

3.2. Nitriding Efficiency and Nitrogen Distribution under Different Process Conditions

Figure 4 presents the SEM microstructure images of all steels, and provides EDS line scanning from the surface to the core, detailing the variation in Cr and N content across the steel specimens. Moreover, the average contents in the interior were measured and indicated below each figure. Figure 4a,b depict the EDS line scanning results of the cross-sections of the 750Ar and 950Ar steels, respectively. The synchronous fluctuation of Cr and N on the sample surface indicates the presence of a significant Cr₂N phase at the surface. As previously analyzed, before the conversion to the Ar atmosphere, partial austenite had undergone sequential decomposition due to excessive nitrogen infiltration, resulting in a mixture of austenite, α-Fe, and Cr₂N phases at the grain boundaries. Following the atmosphere switch, nitrogen escaped from the austenite, leading to a reduction in nitrogen content and ultimately ceasing the decomposition process. Although surface layers experience faster nitrogen infiltration and escape due to their proximity to the free surface, the nitrogen content in both surface and core areas is similar, indicating the potential of the new process to achieve effective and precise control over the nitriding process.

The 1150Ar steel underwent an atmosphere switch before austenite decomposition, preventing the generation of α-Fe and Cr₂N phases, resulting in a uniform austenitic structure with sustained high nitrogen content, as illustrated in Figure 4c. In contrast, the 1320Ar steel prematurely switched from N₂ atmosphere to Ar, leading to inadequate infiltration and excessive escape of nitrogen, causing a gradual decrease in solid-soluble nitrogen in the high-temperature δ-Fe and austenite, which eventually transformed into α-Fe with relatively low nitrogen content during cooling, as shown in Figure 4d.

Figure 4e demonstrates the SEM line scanning result along the cross-section of the 1320N2 steel, highlighting the presence of a layer rich in N and Cr elements, indicating the existence of Cr₂N. The thickness of this layer is approximately 150 μm. Following the solution treatment, these Cr₂N phases were redissolved, eventually forming a single nitrogen-enriched austenitic phase, as shown in Figure 4f.

Figure 5 depicts the hardness distribution across the cross-sections of the steels under different atmosphere-switching conditions. The results reveal higher surface hardness values for the 750Ar, 950Ar, and 1150Ar steels compared to the cores. Furthermore, lower switching temperatures correspond to higher surface hardness values for the respective steel specimens, indicating a higher presence of α-Fe and Cr₂N phases in the surface layers, consistent with the analysis results from Figure 4. In terms of core hardness, the 750Ar steel exhibits the highest average hardness, approximately 340 HV, while the hardness of the 950Ar and 1150Ar steels are comparable, around 300 HV. Surprisingly, the surface hardness of the 1320Ar steel is equivalent to its interior, attributed to its conversion to an Ar atmosphere at the onset of cooling. The nitrogen escaped from the high-temperature δ-Fe and austenite during cooling in the Ar atmosphere, which eventually leads to the formation of α-Fe, without a nitrogen-enriched zone in the surface layer. Additionally, α-Fe has higher strength compared to austenite; hence, despite its much lower nitrogen content compared to the other three steels, the core hardness is highest, similar to that of the 750Ar steel.

3.3. Comparative Performance Advantages of SNAT Processing

Figure 6 presents the engineering stress–strain curves of the steels, with their tensile properties summarized in

Table 4. Tensile mechanical properties of the as-prepared steels.

Sample Code	σb (MPa)	σ0.2 (MPa)	Elongation (%)
750Ar	733 ± 12	437 ± 9	21.8 ± 2.1
950Ar	674 ± 15	411 ± 9	22.4 ± 1.4
1150Ar	1030 ± 20	749 ± 33	38.7 ± 3.7
1320Ar	824 ± 24	483 ± 10	31.9 ± 2.0
1320N2	843 ± 28	578 ± 26	30.2 ± 3.2
1320N2-ST	968 ± 31	676 ± 20	39.5 ± 3.5

. Figure 6a compares the tensile mechanical properties at different transition temperatures. Among these, the 1150Ar steel demonstrates the most outstanding strength and ductility, with a yield strength of 749 MPa, ultimate tensile strength of 1030 MPa, and elongation of 38.7%. This indicates that high-nitrogen austenitic steel can achieve the best strength–ductility combination. The mechanical properties of the 1320Ar steel, with a single-phase ferrite matrix, are slightly inferior due to its lower nitrogen content compared to the former. Meanwhile, both 750Ar and 950Ar steels exhibit the poorest mechanical properties, with elongation decreasing to approximately 20%, primarily attributed to the generation of Cr₂N at grain boundaries. During tensile testing, the presence of Cr₂N precipitates interacts with dislocations, resulting in the formation of dislocation tangles around the nitrides. Consequently, the Cr₂N precipitates act as crack initiation sites, accelerating crack propagation, thus reducing tensile strength and deteriorating ductility. This phenomenon has also been reported by Hu et al. [30] and Zheng et al. [31].

Figure 6b compares the mechanical properties of steel processed through three different routes. The 1320N2 steel exhibits the poorest strength and ductility, which is attributed to the significant formation of nitrides along grain boundaries. However, after undergoing solution treatment, the nitride precipitates in 1320N2 steel were redissolved into the austenite phase in 1320N2-ST steel, leading to a significant enhancement in tensile properties, including a yield strength of 676 MPa, a tensile strength of 968 MPa, and a fracture strain of 39.5%. This improvement can be attributed to the solid solution strengthening provided by nitrogen in solid solution. On the other hand, 1150Ar steel exhibits the highest strength, along with ductility comparable to that of 1320N2-ST. In all, the efficiency and cost-effectiveness advantages of 1150Ar steel are apparent.

Figure 7 illustrates the central area morphology of tensile fracture of steels subjected to various process parameters. Across all tested specimens, alongside the unavoidable small voids inherent in powder metallurgy products, a predominant tough fracture characteristic with dimples and dimple plateaus was observed, indicating they follow the typical mechanism of ductile fracture associated with pore coalescence and growth. Particularly, Figure 7e reveals a distinct laminar fracture morphology in 1320N2 steel, potentially linked to the presence of a ferrite matrix with nitrides. Conversely, this characteristic was absent in both 1150Ar and 1320N2-ST specimens, showcasing instead numerous deep dimples and thereby demonstrating exceptional fracture toughness. This superior performance can primarily be ascribed to their homogeneous austenitic matrix and the absence of chromium nitride.

The electrochemical measurements depicted in Figure 8 and summarized in

Table 5. Electrochemical performance obtained in 3.5 wt.% NaCl solution for as-prepared steels.

Sample Code	Ecorr (V)	Icorr (mA/cm2)	Rp (kΩ)	C. Rate (mm/y)
750Ar	−0.46	0.76	5.3	13
950Ar	−0.42	0.24	10.5	4
1150Ar	−0.21	0.06	42.1	1.1
1320Ar	−0.29	0.13	19.6	2
1320N2	−0.43	7.13	15.2	82
1320N2-ST	−0.24	0.10	30.3	1.4

reveal notable differences among the steels. Figure 8a illustrates the potentiodynamic polarization curves of the 750Ar, 950Ar, 1150Ar, and 1320Ar steels. It is evident that all four steels demonstrate similar activation–passivation behaviors. The polarization curves encompass activation, transition, passivation–activation, and activation regions. Detailed in Table 5, the corrosion potentials of the 750Ar, 950Ar, 1150Ar, and 1320Ar steels are −0.46 V, −0.42 V, −0.21 V, and −0.29 V, respectively. The corrosion current density values of 0.76 mA/cm², 0.24 mA/cm², 0.06 mA/cm², and 0.13 mA/cm², along with the corresponding annual corrosion rates of 13 mm/y, 4 mm/y, 1.1 mm/y, and 2 mm/y, respectively. Of note, the 750Ar and 950Ar steels, comprising austenite, α-Fe, and Cr₂N, foster the formation of numerous galvanic cells among these phases, leading to accelerated corrosion rates. Consequently, they exhibit the highest corrosion current density and annual corrosion rate, indicating the least corrosion resistance among the tested steels. Although the 1320Ar steel features an α-Fe matrix, its lack of nitrides contributes to better corrosion resistance compared to the former. Particularly noteworthy is the 1150Ar steel, featuring a single austenitic phase, which manifests the highest corrosion potential and the lowest corrosion current density and annual corrosion rate.

Figure 8b illustrates that both the 1320N2-ST and 1150Ar steels exhibit higher polarization resistance and lower corrosion current density compared to 1320N2 steel. This improvement can be attributed to the high nitrogen content in the austenitic matrix [32]. However, the microstructure of 1320N2 includes austenite, α-Fe, and Cr₂N. On one hand, the precipitation of Cr₂N results in the depletion of Cr and N in austenite, which will reduce the corrosion resistance. On the other hand, the nitrides and α-Fe increase the number of galvanic cells in the electrochemical corrosion system and aggravate the intergranular corrosion. Therefore, by strictly controlling the N content and reducing the precipitation of nitrides, the corrosion resistance of the high-nitrogen austenitic stainless steel can be effectively improved.

Figure 9 compares mechanical properties, corrosion resistance, and material costs among stainless steels produced via the MIM process. The results indicate that the material under investigation demonstrates promising application potential and commercial value when compared to traditional types such as 316L [33], 17-4PH [34], and 420 stainless steel [35]. The 1150Ar steel demonstrates outstanding mechanical properties and corrosion resistance while remaining cost-competitive. Moreover, the incorporation of atmosphere switching in the production process significantly reduces manufacturing time and costs, highlighting its potential for broad industrial utilization.

4. Conclusions

In this work, we proposed a novel method to achieve a homogeneous microstructure of single-phase austenite with excellent properties in MIM high-nitrogen austenitic stainless steel. This method eliminates the need for traditional solution treatment by switching the cooling atmosphere from N₂ to Ar at an appropriate temperature. Based on a comprehensive analysis of the microstructure and performance of steels subjected to N₂ cooling, N₂-Ar cooling, and solution treatment, the following conclusions are drawn:

(1)

Cooling in a N₂ atmosphere initially transforms the matrix to a nitrogen-rich austenite. Excessive nitrogen infiltration precipitates Cr₂N, inducing partial austenite decomposition and ultimately forming a multiphase structure of austenite, α-Fe, and Cr₂N.

(2)

Switching atmosphere from N₂ to Ar reverses the nitriding reaction, and the earlier the switch, the shorter the nitriding time and the longer the denitriding time. A balance between these two processes is needed to achieve a high-nitrogen, chemically uniform single-phase austenitic structure.

(3)

Optimal properties were observed at a switching temperature of 1150 °C, where the steel exhibited a yield strength of 749 MPa, ultimate tensile strength of 1030 MPa, elongation of 38.7%, and corrosion current density of 0.06 mA/cm². These properties surpass those of N₂-cooled and solution-treated steels, as well as other MIM stainless steels.