Isothermal aging treatment and its effect on the mechanical properties of Fe-Cr-Mn high-nitrogen austenitic steel

The high-temperature solution nitriding process is a suitable treatment for producing high-nitrogen steels by the diffusion of nitrogen from the surface to the center of steels under high nitrogen pressure at high temperatures. On the other hand, long-term solution nitriding at high temperatures can cause the formation of coarse-grained austenite. This study focused on the grain size, strength, and elongation properties of Fe-Cr-Mn high-nitrogen austenitic steels by isothermal aging. For this purpose, we produced high-nitrogen steels by three-step phase transformations: (i) high-temperature solution nitriding, (ii) isothermal aging, and (iii) reaustenitization. After solution nitriding, high-nitrogen austenitic steel was treated with isothermal aging to induce austenite decomposition. Supersaturated austenite (ϒ) transformed to less nitrogen-saturated austenite (ϒ′) and discontinuous cellular precipitation (DCP) during the isothermal aging treatment. Subsequently, the decomposed structure was reversed into austenite through reaustenitization. The results suggested that austenite grain refinement and increasing hardness were achieved by nucleating reversed austenite grains after reaustenitization treatment. On the other hand, a noticeable increase in elongation appeared after reaustenitization treatment for 30 min.


Introduction
High-nitrogen austenitic stainless steels, with their unique combination of mechanical strength and corrosion resistance, have significant practical implications.Nitrogen, as an austenite stabilizer, plays a crucial role in reducing the formation of ferrite and martensite upon solidification.Understanding the effects of weathering on these steels can lead to improved hardening mechanisms such as solid solution hardening, fine-grain strengthening, and cold working.This research aims to contribute to this understanding, potentially leading to the development of stronger and more corrosion-resistant steels [1][2][3][4][5][6][7][8].
Research has investigated the influence of the addition of nitrogen on the mechanical properties of stainless steel.It was reported that higher mechanical properties could be obtained in HNS [9][10][11][12][13].Despite the superior mechanical properties of HNS, its production presents significant challenges.HNSs can be produced by two methods based on the liquid or solid state of steels: (i) liquid-state treatment, which involves high-pressure electro-slag remelting of steel in a nitrogen gas atmosphere, and (ii) solid-state treatment, which involves a hightemperature solution nitriding process (HTSN) via nitrogen gas absorption on the steel surface.The liquid-state treatment of HNSs is applied with high pressure because of the low solubility of nitrogen in liquids.This production method has high production costs and other limitations, such as an increased risk of crack formation during the hot forming process [2,[14][15][16].In the solute state approach, HNS is produced by long-term HTSN in a nitrogen gas atmosphere at 1000 °C-1200 °C [1,2,[17][18][19][20].The HTSN process allows the dissolution of high nitrogen contents (∼0.5-1.0 wt%) in austenite [21][22][23][24][25][26][27].
One of the significant issues of the HTSN process is that austenite has coarse grain sizes after long-term annealing at high temperatures.To solve this issue, a range of alternative approaches can be considered based on obtaining finer grains via austenite recrystallization from cold-deformed structures or austenite reversion from precipitates.Cold deformation before subsequent recrystallization annealing is a general approach for obtaining fine-grained structures of austenitic HNS.Li et al [28] reported that cold rolling and subsequent annealing produced fine-grained structures in nickel-free HNS, in which the strength and hardness increased, and the ductility decreased with decreasing grain size.According to another study by Talha et al [29], cold-deformed Nifree HNSs highlighted the impact of nitrogen, as indicated by increased macrohardness with increased nitrogen content and significant reductions in thickness.These authors noted that no martensite formed during cold working.However, cold deformation accelerates the formation of nitrides and intermetallics.There is a complex relationship between precipitation and recrystallization in HNS.Shi et al [30] reported that the precipitation of second-phase particles induced recrystallization aging at 750 °C.The second-phase particles nucleated at dislocations, grain boundaries, and subgrain boundaries, hindering the formation of recrystallization nuclei.
Apart from the grain refinement achieved through cold deformation, an alternative method involves the formation of precipitates to initiate the nucleation of new grains within these precipitates.Nakada et al [31] studied the grain refinement of nickel-free HNS (Fe-25Cr-1 N mass% alloy) by a two-step phase transformation process, including a ferritic initial structure.Initially, the austenitic steel underwent isothermal heat treatment at 927 °C, forming a lamellar structure composed of ferrite and Cr 2 N. Subsequently, the specimen with this eutectoid structure was reverted to austenite through heating.Therefore, the reversed austenite grains had many nucleation sites at the grain and eutectoid block boundaries.Their results suggested that austenite grain refinement can be obtained by isothermal aging.Onomoto et al [32] examined the effect of grain refinement by isothermal aging and reversion heat treatment methods on the tensile properties of Fe-25Cr-1N alloy.The results showed that as the grain size decreased, the elongation increased significantly.Mohammadzadeh et al [33] applied a similar technique and reported that eutectoid transformation from austenite to ferrite and Cr 2 N occurred during isothermal heat treatment below 1000 °C.According to their results, Cr 2 N does not dissolve entirely during reaustenitization at 1200 °C for one hour.An intermediate homogenizing heat treatment cycle at 1200 °C for 10 h was applied to address this before the grain refinement process.After long-term homogenization treatment, the average grain size reached 100 μm.Reaustenitizing for 10 min after homogenization resulted in a fine-grained austenitic structure with an average grain size of approximately 20 μm.
The formation of nitrides in HNS has been studied extensively.However, these studies mainly address the precipitation of nitrides during thermal processes such as hot forming, heat treatment, service environments at high temperatures, and welding [34][35][36][37][38][39][40].Nitride precipitation usually occurs between 700 and 1000 °C, depending on the chemical composition of the HNS alloy.The formation of nitride precipitates at these temperatures is studied via isothermal aging treatment.Shi et al [34] examined the morphology and precipitation kinetics of Cr 2 N after hot deformation in the isothermal aging process of Fe-18Cr-12Mn-0.48N HNS.These authors indicated that the nose temperature during nitride precipitation is approximately 800 °C and that Cr 2 N precipitation transforms from an initial granular to lamellar form with increasing aging time.Qin et al [36] reported similar results for Cr 2 N precipitation in a Fe-21Cr-19Mn-0.69N HNS alloy.However, the sigma phase can also be found in HNS.In a study of the aging precipitation behavior of hot rolled Fe-25Cr-3Mn-7Mo-0.5 N HNS, Zhang et al [35] detected sigma precipitates that formed as intergranular, cellular, and intragranular.These authors indicated that the sigma phase transformed into Cr 2 N and the R phase through Cr and Mo diffusion at 700 °C-1100 °C.According to their more recent study of the same material, Zhang et al [37] reported that increasing nitrogen promotes the formation of intergranular precipitates.Thus, grain refinement of austenite can occur.Li et al [38] investigated the aging precipitation behavior of 18Cr-16Mn-2Mo-1.1 N highnitrogen austenitic stainless steel and its influence on its mechanical properties.They found that cellular Cr 2 N increased the hardness but reduced the ultimate tensile strength, yield strength, and elongation.Wang et al [39] researched the enhancing yield strength of 18Mn18Cr0.6N high-nitrogen austenitic stainless steel from Fe-18Cr-15Mn-0.66N stainless steel.They observed that the yield strength increased due to the precipitationstrengthening effect of Cr 2 N. Dandekar et al [40] investigated the effect of isothermal aging on the microstructure and mechanical properties of UNS S32101 lean duplex stainless steel.They suggested that ductility decreases, and yield/ultimate tensile strength increases with aging time.
Research on the HTSN process has centered primarily on surface treatment of duplex stainless steels rather than manufacturing HNS [41].The manufacture of austenitic steels by HTSNs can depend on their initial structure, such as ferrite or ferrite+austenite.As mentioned above, the aging precipitation behavior of highnitrogen austenitic steels and its effect were investigated.However, no previous study has investigated reaustenitization treatment via HTSN after isothermal aging of high-manganese austenitic HNSs using a duplex initial microstructure.This paper examines isothermal aging and reaustenitization treatment and its effect on the mechanical properties of an austenitic HNS alloy.We produced the samples in three steps: (i) hightemperature solution nitriding, (ii) isothermal aging, and (iii) reaustenitization.After hot forging-hot rolling, and cold rolling, a duplex microstructure of ferrite and austenite was obtained.The duplex structure was transformed into a fully austenitic microstructure by HTSN at 1200 °C.Nitrogen-supersaturated austenite (ϒ) decomposed to a less nitrogen-saturated matrix (ϒ′) and discontinuous cellular precipitates during isothermal aging.Finally, reversed austenite grains were obtained at the reaustenitization stage.The present study aimed to examine austenite grain refinement via precipitates and investigate its effect on the mechanical properties of austenitic HNS.Microstructural characterization was carried out by light microscopy, scanning electron microscopy, and x-ray diffraction methods.The mechanical properties were determined by observing the hardness, ultimate tensile strength, and elongation via microhardness and tensile tests.

Experimental approach
This study produced a newly developed duplex steel, Fe-20Cr-14Mn-2Ni-3Mo alloy, as the base material for high-temperature solution nitriding (HTSN).The chemical composition of the duplex steel is listed in table 1.A 30 kg ingot (∅110 mm × 500 mm) was produced by induction melting in a vacuum and then cast into a metallic mold.The ingot was homogenized by annealing at 1200 °C for 48 h and then cooled to room temperature at 40 K min −1 .The ingot was hot forged into a 26 mm thick plate after being reheated and held at 1200 °C for 1 h.The hot forged plate was machined to a thickness of 24 mm to remove surface errors before hot rolling.After reheating at 1200 °C for 1 h, the plate was hot rolled at 900 °C-1150 °C to a final thickness of 4 mm through 12 passes.The hot-rolled sheet was machined to 3.5 mm to remove the scale.Subsequently, the hot-rolled sheet was cold rolled with 88% cold reduction to a thickness of 0.4 mm through 20 passes.Then, the cold-rolled plates were cut into 10 * 10 mm specimens.All the specimens were subjected to three treatments: HTSN treatment, isothermal aging, and reaustenitization.
Before the heat treatments, thermodynamic calculations were performed.Figure 1 shows the calculated phase fractions for the duplex steel using ThermoCalc software and the TCFE13 database.According to the equilibrium calculations in figure 1, solidification was completed with the δ-ferrite phase.In addition to the δferrite phase, austenite, M 23 C 6 , σ (sigma), and Laves phases are predicted.The solidification process from the liquid (L) phase to 600 °C can be defined as The thermodynamic calculations show that the alloy has a single-phase region of δ-ferrite at temperatures ranging from 1080 °C to 1390 °C, allowing for the appropriate HTSN temperature adjustment.
In figure 2, a schematic diagram of the heat treatment routes used for (i) high-temperature solution nitriding (HTSN), (ii) isothermal aging, and (iii) reaustenitization is given.The details of the heat treatments are as follows: (i) The HTSN was applied to 0.4 mm thick 10 * 10 mm 2 cold rolled duplex steel sheets at 1200 °C in a highpurity nitrogen atmosphere at a pressure of 2.5 bar, with varying durations between 15 and 120 min.After the HTSN, the samples cooled to room temperature at 200 K/s.The ferritic+austenitic duplex structure was transformed to a complete austenitic phase at 1200 °C for 30 min by HTSN.
(ii) An isothermal aging treatment was applied to the 30 min of HTSN-treated samples to obtain austenite grain refinement via the ART annealing approach.The HTSN-treated samples were isothermally aged at temperatures ranging from 700 °C-1000 °C for 20 min under a high-purity argon atmosphere and then water quenched to room temperature.The microstructure decomposed during the isothermal aging from supersaturated austenite to saturated austenite and chromium nitride precipitates.
(iii) After the isothermal aging, reaustenitization was applied for the reversion treatment, and reversed austenite grains nucleated.The reaustenitization step was applied to the isothermally aged sample 900 °C.The reaustenitization sample was annealed at 1200 °C for 30 min under a high-purity argon atmosphere.
The nitrogen content in each specimen was determined via LECO analysis.Light microscopy (LM), scanning electron microscopy (SEM), energy dispersive x-ray spectroscopy (EDS), and x-ray diffraction (XRD) patterns were used to observe the microstructure and morphology of the samples after etching with aqua regia, 10% aqueous oxalic acid, and Beraha ll solutions.The hardness and tensile strength were determined from the mechanical properties.The Micro-Vickers hardness was measured on a FutureTech FV700 test machine at five kgf.Uniform tensile tests were performed at room temperature on the Instron universal test machine.The tensile subsized specimens with the geometry indicated in figure 3 were machined according to ASTM E8/DIN EN ISO 6892-1.

Results and discussion
3.1.Microstructural characterization 3.1.1.Microstructure of the as-cast specimen Figure 4 shows the microstructures of the as-cast sample.The light microscope image in figure 4(a) reveals that a typical duplex steel microstructure of austenite and δ-ferrite was obtained.It should be noted that Widmanstätten austenite appears along the delta ferrite grain boundaries and grows further inside the grains.This morphology can be described by the slow cooling rate during solidification.According to a study of duplex stainless steels by Ohmori et al [42], Widmanstätten austenite laths form at grain boundaries of delta ferrite or nucleate directly in the delta ferrite grains.Precipitation can be found in a high temperature range between 700 °C and 1100 °C.Wu et al [43][44][45] also investigated the formation, morphology, and effect of Widmanstätten austenite on the fatigue behavior of duplex stainless steel.It was found that Widmanstatten austenite formed during slow cooling from high temperature deteriorated the impact toughness and the tensile elongation.The elemental compositions of the austenite and δ-ferrite phases were measured via SEM-EDS analysis and are summarized in figure 4(b).The ferritic matrix is rich in chromium and molybdenum, while the manganese and nickel contents are more significant in austenite.

Microstructures after homogenization treatment, hot rolling, and cold rolling
A homogenization treatment was applied to the ingot for 48 h at 1200 °C to dissolve Widmanstätten austenite in the ferritic matrix.After the homogenization treatment, Widmanstätten austenite transformed to intergranular and intragranular austenite, as shown from the microstructure of the ingot center in figure 5(a).The SEM image and EDS analysis are shown in figure 5(b).The intermetallic phases sigma (σ) and chi (χ) were present in the intragranular austenite areas.Escriba et al [46] reported that sigma usually nucleates and grows within intragranular austenite grains.Chi precipitates with a more spherical morphology within ferrite at the boundaries of ferrite and intragranular austenite.The sigma usually precipitates a lamellar morphology due to the eutectoid transformation of ferrite to the sigma and austenite.However, the sigma formation can lack eutectoid-like morphology as reported by Inacio et al [47].As observed in the present study, the sigma phase was not lamellar morphology and formed separately within austenite/austenite interfaces.Recent studies have shown that the sigma phase and chi phase precipitate at high Cr-concentrated regions.[47][48][49][50][51]. EDS analysis revealed that sigma and chi contain high levels of Mo and Cr.On the other hand, these intermetallic phases are richer in Mo than ferrite.It was concluded that intermetallic phases were present due to the cooling of the ingot at a slow cooling rate of 40 K min −1 after homogenization annealing.Figure 6 shows the XRD patterns taken from the surface of the as-cast ingot center and after homogenization annealing.XRD analysis revealed that the sample consisted of ferrite and austenite phases.The (211) α and (111) γ peaks were observed in the sample from the as-cast ingot.However, sharp (110) α and (220) γ peaks were more intense after homogenization annealing, indicating texture formation in both phases [52].On the other hand, the intermetallic phases sigma and chi shown in figure 5 were not found via XRD analysis due to their trace amounts.
The homogenized annealed ingot was hot forged and hot rolled.The final microstructure of the hot-rolled duplex steel is shown in figure 7(b).During hot deformation, the austenite islands become aligned in the hotrolling direction.The intermetallic phases were not observed due to reheating and rapid cooling after hot rolling.The volume fractions of the duplex microstructure include approximately 30 vol.-% austenite and 70 vol.-%ferrite.
After hot rolling, the hot-rolled plate was cold rolled to obtain thin-sheet samples for HTSN experiments.Figure 7(b) shows that the microstructure of the homogenized sample changes from equiaxed grains to more laminated microstructures after hot and cold rolling.

High-temperature solution nitriding
High-temperature solution nitriding (HTSN) at 1200 °C for different durations between 15 and 120 min was applied to cold rolled duplex steel sheets with a thickness of 0.4 mm to obtain a complete austenite microstructure.Light micrographs of the specimen surfaces before and after HTSN are shown in figure 8. Nitrogen diffusion during the HTSN treatment after a 30 min duration was achieved through exposure to a high-purity N 2 atmosphere, resulting in the evolution of a fully austenitic microstructure.However, the ferritic matrix and the austenite phase still appear in the sample's microstructure after HTSN treatment for 15 min.
The presence of ferrite in the HTNS sample after 15 min was also confirmed by x-ray diffraction (XRD), as shown in figure 9.However, the intensity of the ferrite peak decreased compared to that of the sample before HTSN.XRD analysis also revealed that no ferrite peak was observed in the samples obtained by HTSN treatment for durations ranging from 30 min to 120 min.Cross-sectional examinations were also carried out to confirm that the microstructure was utterly transformed into austenite.Figure 10 shows cross-sectional SEM images of the samples.After HTSN for 30 min, the microstructure was fully austenitic.SEM images show that the austenite grains are twinned.
The average nitrogen contents of the samples after HTSN were measured using LECO analysis.The nitrogen contents in the samples after HTSN for 30 min, 60 min, and 120 min were 1.60%, 1.73%, and 1.68%, respectively.According to the empirical formula (PREN 16 = %Cr + 3.3%Mo + 16%N) by Merello et al [53], the steel's PREN value was 57 after adding 1.60 wt% nitrogen via the HTSN.

Post-heat treatments after the HTSN 3.3.1. Isothermal aging treatment
After the HTSN at 1200 °C, the average austenite grain size was approximately 52.3 μm.An isothermal aging approach was applied to obtain a finer austenite grain size.Before isothermal aging treatment, a thermodynamic calculation was performed to determine the proper temperature range.Figure 11 shows the equilibrium phase fractions as a function of temperature for the chemical composition after the HTSN.According to the calculation results, secondary precipitates of phases such as Cr 2 N, sigma, Laves, and M 23 C 6 are predicted.The mass fractions of the M 23 C 6 , sigma, and Laves phases are lower than 3 vol-%.Cr 2 N is the main stable phase over the temperature range from 700 to 1100 °C.Thus, a temperature range between 700 and 1000 °C was chosen for the isothermal aging experiments.To vary the phase fractions of Cr 2 N, isothermal aging was performed at four temperatures at increments of 100 °C between 700 and 1000 °C for 20 min.
SEM images of the isothermally aged samples are shown in figure 12. EDS analysis was also performed to determine the chemical composition of the precipitates.SEM-EDS point analysis results of isothermally aged samples are given in table 2. It was found that only austenite and chromium nitride phases were present in the microstructure.The chromium nitride phase has higher Cr and N contents than austenite.Figure 12(a) shows that the chromium nitride nucleated at the austenite grain boundaries at low temperatures.At temperatures above 900 °C, complete precipitation of chromium nitride is observed.According to the calculation results of ThermoCalc software, these precipitates were defined as Cr 2 N. Knutsen et al [54] reported that chromium nitrides form discontinuously in high-nitrogen Cr-Mn austenitic steels upon aging at 700 °C-1000 °C.The decomposition of a supersaturated solid solution (ϒ) into a less nitrogen-saturated matrix phase (ϒ′) and chromium nitride (Cr 2 N) at a moving boundary is known as discontinuous cellular precipitation (DCP).Thus, DCP corresponds to the following reaction: γ→γ′+ Cr 2 N. The chromium nitrides grew discontinuously toward the interior of the austenite grains.During the DCP, chromium nitride lamellae formed.Single-phase supersaturated ϒ transforms into a solute-depleted ϒ′ matrix and grain boundary cells consisting of alternate lamellae of a precipitation phase, similar to the precipitation behavior of pearlite in carbon steels.This precipitation type is called 'nitrogenous pearlite' in high-nitrogen austenitic steels.According to the results of EDS analysis, by increasing the temperature during isothermal aging, the solute-depleted ϒ′ phase has lower amounts of the former ferrite elements Cr and Mo than does the ϒ′ phase, while the Mn amount increases.

Reaustenitization
The isothermal aging condition is the initial microstructure before reaustenitization treatments to obtain finer austenite grains than the samples after HTNS.After isothermal aging at 900 °C, discontinuous precipitates completely covered the sample.Above 900 °C, the mean size of the Cr 2 N lamellae increased with increasing isothermal aging temperature.Therefore, isothermal aging treatment was applied at 900 °C.In comparison, the microstructure of the sample after HTSN is also given in this section.The microstructural changes resulting from reaustenitization are presented below.
Reaustenitization was carried out at 1200 °C for 30 min.Sample A represents the sample under HTNS conditions without post-heat treatment, while the sample after post-heat treatment is denoted Sample B. The sample codes and process parameters are given in table 3.
The SEM images of the samples after HTSN and reaustenitization are given in figure 13.It was observed that reversed austenite grains formed from less nitrogen-saturated austenite and cellular nitride precipitations by reaustenitization precipitates by reaustenitization.
Nakada et al, Onomoto et al and Mohammadzadeh et al [31][32][33] studied isothermal aging and reaustenitization treatments on nickel and manganese-free austenitic stainless steel produced by HTSN from a ferritic Fe-Cr-Mo steel alloy.They reported that grain refinement was achieved by applying this two-stage heat treatment procedure.
Grain size measurements were carried out to investigate the effect of isothermal aging on the final grain size.Figure 14 shows light microscopy images of the samples after HTSN and two-step heat treatment, isothermal aging, and reaustenitization.It can be observed that grain refinement occurs after the two-step heat treatment.The average grain size of Sample A is 52.3±22.2μm, and that of Sample B is 40.1 ± 17.4 μm.
Figure 15 shows the grain size distribution curves of the samples before and after reaustenitization.According to the results, a more homogeneous grain size distribution is obtained after the reaustenitization treatment (Sample B).It is evident that 22% of the grains are more significant than 100 μm after the HTSN treatment (Sample A).It was observed that the austenite grains recrystallized during the isothermal aging treatment.New grains less than 15 μm in length formed during the reaustenitization step.The recrystallization twins are also observed in figure 15(b).

Mechanical properties
The mechanical properties of the samples were determined by performing microhardness and tensile tests at room temperature.The ultimate tensile strength (UTS), total elongation (TE), and product of the UTS and total elongation (PSE) were measured via tensile testing.The microhardness and tensile test results are summarized in table 4. The hardness of the sample after the HTSN (Sample A) reaches approximately 342 HV.The reaustenitization treatment resulted in 348 HV in Sample B, similar to Sample A's hardness.Following the reaustenitization treatment, the total elongation increased to 41% from 29%.Sample B obtained the optimum combination of tensile properties by exhibiting UTS of 910 MPa, a TE of 41%, and a PSE of 37.31 GPa%.The reaustenitization based on ART annealing approach effectively refined grain size without the need for deformation processes.Some coarse grains within the 100-200 μm range were present in Sample A, but these grains were absent in Sample B due to the reaustenitization treatment.Additionally, Sample B includes reversed austenite grains smaller than 15 μm, formed after the isothermal aging and reaustenitization treatments.The absence of coarse grains and the formation of new fine grains decreased the average grain size.The decrease in the average grain size led to an improvement in mechanical properties.The fracture modes of the steels were studied, and SEM micrographs of the fracture surfaces of the tensile tested samples are presented in figure 16.The morphologies of the fracture surfaces show that ductile fracture is the main fracture mode for both samples.The fracture surface of Sample A is at a higher magnification in figure 16(b).shows shear bands, as marked by white arrows.This deformation mode as shear bands indicates the localization of plastic deformation at the necking region.The fracture surface of Sample B is given at a higher magnification in figure 16(d).The dimples' size in Sample B's fracture surface is greater than that of Sample A, indicating that Sample B has more excellent tensile ductility.

Conclusions
This paper investigated the effect of isothermal aging treatment on the microstructure and mechanical properties of Fe-Cr-Mn high-nitrogen austenitic steel.The following conclusions can be drawn from the present study: (1) In the as-cast condition of Fe-20Cr-14Mn-2Ni-3Mo duplex stainless steel, the microstructure revealed the presence of Widmanstatten austenite at the delta ferrite grain boundaries.After the homogenization treatment, the Widmanstatten austenite transformed to intergranular and intragranular austenite in the ferritic matrix.The intermetallic phases sigma (σ) and chi (χ) were also present in the homogenized sample.However, intermetallic precipitates were not detected following the reheating and hot rolling steps.Finally, after solution nitriding at 1200 °C, Fe-20Cr-14Mn-2Ni-3Mo-1.6 N austenitic steel is obtained.
(2) A two-stage heat treatment cycle was applied to the solution-nitrided samples to achieve a finer austenite grain size and distribution by austenite reversion, including isothermal aging and reaustenitization steps.At the isothermal aging step, supersaturated austenite (ϒ) decomposed to saturated austenite (ϒ′), and Cr 2 N was named 'nitrogenous pearlite' at temperatures between 700 and 1000 °C for 20 min.Cr 2 N precipitates began to form on the austenite grain boundaries at 700 °C.Increasing the isothermal aging temperature, more cellular Cr 2 N is found in the austenite matrix.At temperatures above 900 °C, complete precipitation of chromium nitride is observed.Above 900 °C, the mean size of the Cr 2 N lamellae increased with increasing isothermal aging temperature.Therefore, isothermal aging treatment was applied at 900 °C.
(3) The isothermally aged sample was reaustenitized at 1200 °C for 30 min.The microstructure of the isothermally aged sample completely transformed into the supersaturated austenite phase after reaustenitization, and the nucleation of reversed austenite grains provided a fine-grained austenitic microstructure.The average austenite grain size reduced to 40 μm from 52 μm.(4) The improvement in mechanical properties aligns with the refinement of the austenite grain size.It was observed that new small grains are formed after the isothermal aging and reaustenitization treatments.Additionally, the presence of coarse grains within the range of 100-200 microns, which existed following the HTSN process, were not seen after the reaustenitization treatment.Higher TE and the UTS * TE product values indicate more excellent ductility, while a slight decrease in UTS is observed.According to the investigation of the fracture surface of the tensile samples, larger dimples implying more excellent ductility were found in the fracture surface of the sample with reversed austenite.

Figure 2 .
Figure 2. Schematic diagram of the heat treatment routes.

Figure 3 .
Figure 3.The scheme of the tensile specimen.

Figure 4 .
Figure 4. Microstructures of the as-cast ingot center; (a) LM micrograph; (b) SEM micrograph and SEM-EDS point analysis results (chemically etched with Beraha ll).

Figure 5 .
Figure 5. (a) LM micrograph and (b) SEM micrograph and SEM-EDS point analysis results after homogenization annealing for 48 h at 1200 °C (chemically etched with Beraha ll).

Figure 6 .
Figure 6.XRD results for samples (a) as-cast and (b) after homogenization annealing.

Figure 7 .
Figure 7.Light micrographs of the samples after (a) hot rolling and (b) cold rolling (chemically etched with Beraha ll).

Figure 9 .
Figure 9. XRD patterns of the samples before and after the HTSN.

Figure 14 .
Figure 14.Light micrographs of (a) Sample A and (b) Sample B (chemically etched with aqua regia).

Figure 15 .
Figure 15.Grain size distribution curves of HTSN-treated Sample A and reaustenitized Sample B after the HTSN.

Figure 16 .
Figure 16.Scanning electron micrographs of fracture surfaces of tensile-tested samples (a), (b) Sample A and (c), (d) Sample B.

Table 1 .
Chemical composition of stainless steel.

Table 2 .
SEM-EDS point analysis results of isothermally aged samples.

Table 3 .
Experimental sample codes and process parameters.

Table 4 .
Microhardness and tensile test results.