20Cr25NiNbRE austenitic stainless steel, and preparation method and application thereof

Abstract

The present application belongs to the field of stainless steel material and processing technology for ultra-supercritical boiler and high-temperature nuclear power reactor internals, and relates to a 20Cr25NiNb austenitic stainless steel, a preparation method and application thereof. The composition of the austenitic stainless steel comprises C, Cr, Ni, Mn, Si, Mo, Nb, N, B, Ce, Al, Ti, P, S, O, H and Fe. The present application adopts vacuum pressurized smelting and nitrogen control technology, and adds a multi-stage thermal mechanical treatment path of'short-time annealing + secondary three-up-three-down'. Through the synergistic effect of component purification, Ce functionalization regulation and innovative thermal processing technology, the present application effectively overcomes the high-temperature hot brittleness and unstable structure problems of traditional 20Cr25NiNb austenitic stainless steel, while maintaining high strength and oxidation resistance, significantly improving high-temperature hot ductility and long-term service reliability.

Description

Technical Field

[0001] This invention belongs to the field of stainless steel materials and processing technology for internal components of ultra-supercritical boilers and high-temperature nuclear power reactors, and relates to a 20Cr25NiNbRE austenitic stainless steel, its preparation method and application. Background Technology

[0002] Austenitic stainless steel, due to its excellent high-temperature strength, good creep resistance, and outstanding oxidation and corrosion resistance, is widely used in high-temperature and extreme service environments such as ultra-supercritical thermal power units and advanced nuclear reactor internals. Under such conditions, the material is subjected to the combined effects of high temperatures above approximately 600°C, complex stress, and corrosive media for extended periods. This requires not only high high-temperature strength and long service life, but also good high-temperature thermal ductility to meet the demands of high-plasticity hot forming, and the ability to alleviate stress concentration and delay creep damage and crack initiation through plastic deformation during service.

[0003] 20Cr25NiNbRE austenitic stainless steel, with properties similar to the existing NF709 / Alloy709 (20Cr25NiNb), is a representative material developed for the aforementioned applications. It achieves a stable austenitic structure through a high Cr and high Ni composition design and relies on MX precipitates such as Nb(C,N) for high-temperature strengthening. However, existing 20Cr25NiNb steel still has significant shortcomings in engineering applications: on the one hand, the material is prone to grain boundary weakening during hot working and high-temperature service, mainly due to the segregation of impurity elements (S, P, O), the presence of coarse inclusions, and the continuous precipitation of MX phases at grain boundaries. 23 Brittle phases such as C6 lead to insufficient high-temperature thermal ductility, increasing the risk of forging cracks and reducing service reliability. On the other hand, the dendritic segregation and uneven distribution of primary precipitates that are common in the as-cast microstructure are difficult to completely eliminate through conventional homogenization, which in turn causes instability in mechanical properties and corrosion resistance, restricting the application of materials in large critical components.

[0004] To address the aforementioned issues, there is an urgent need to develop an austenitic stainless steel that combines high compositional purity with precise microstructure control. This would meet the requirements of next-generation ultra-supercritical and advanced nuclear energy systems for high-performance heat-resistant stainless steel materials. Summary of the Invention

[0005] To address the technical problems existing in the prior art, this invention provides a 20Cr25NiNbRE austenitic stainless steel, its preparation method, and its applications. This steel grade incorporates trace amounts of cerium (Ce, 0.005-0.12 wt%) for microalloying design based on the traditional alloy system. Combined with ultra-high purity smelting and vacuum pressure nitrogen alloying technology, the content of harmful impurities such as oxygen and sulfur, as well as residual elements such as arsenic and tin, is strictly controlled to extremely low levels, thereby improving material purity from the source and refining and optimizing the morphology of inclusions. Simultaneously, addressing the issues of segregation in the as-cast structure and insufficient high-temperature thermal ductility, a multi-stage hot working process centered on "large strain hot deformation-diffusion synergistic homogenization" is constructed. This process sequentially includes ingot homogenization, "three upsetting and three drawing" large deformation forging in the single-phase austenite region, intermediate annealing, secondary upsetting and drawing deformation, and solution and stabilization treatment. Through the synergistic effect of composition regulation and process optimization, not only is the dendritic segregation in the as-cast state effectively weakened and the microstructure homogenization promoted, but grain refinement and second-phase regulation are also achieved. While ensuring the high strength and excellent oxidation resistance of the material, its high-temperature thermal ductility and microstructure stability are significantly improved, thereby effectively improving the hot brittleness problem of traditional similar steel grades during hot working and service.

[0006] The technical solution of this invention is: A type of 20Cr25NiNbRE austenitic stainless steel, with the following chemical composition by mass percentage: C ≤ 0.10%, Cr: 19-21%, Ni: 24-26%, Mn: 0.3-1.5%, Si: 0.1-0.5%, Mo: 1-2%, Nb ≤ 0.30%, N: 0.10-0.30%, B: 0.001-0.003%, Ce: 0.005-0.12%, Al ≤ 0.05%, Ti ≤ 0.02%, P ≤ 0.03%, S ≤ 0.02%, O ≤ 0.003%, H ≤ 0.0002%, with the balance being Fe and unavoidable impurities.

[0007] The chemical composition, by mass percentage, further satisfies the following: Nb: 0.15-0.25%, N: 0.15-0.25%, C: 0.03-0.07%, Ti ≤ 0.01%.

[0008] The chemical composition, by mass percentage, further satisfies the following: P ≤ 0.01%, S ≤ 0.005%, O ≤ 0.002%, H ≤ 0.0002%.

[0009] The chemical composition, by mass percentage, further satisfies: Ce: 0.02-0.08%.

[0010] A method for preparing 20Cr25NiNbRE austenitic stainless steel includes the following steps: (1) Smelting, nitrogen alloying and rare earth treatment: Add the required raw materials other than rare earth metal Ce to the vacuum pressure induction furnace according to the target composition. After refining at 1550-1600℃ for 10-15 minutes, nitrogen gas is introduced into the furnace for nitrogen alloying. Then, rare earth metal Ce is added under nitrogen protection for rare earth treatment to obtain molten steel to be poured. (2) Casting and homogenization treatment: molten steel is cast at 1500-1540℃ with nitrogen pressure maintained at 0.18-0.25 MPa; after casting, the pressure is released to atmospheric pressure, the furnace is cooled to 900-950℃ and then demolded at high temperature, followed by water cooling to room temperature to obtain ingots; the ingots are heated to 1180-1250℃ for homogenization treatment and held for 4-6 h. (3) First large plastic deformation: The ingot is subjected to the first "three upsetting and three drawing" in the single-phase austenite temperature range of 1180-1250℃, and the cumulative equivalent deformation is ≥60%; (4) Intermediate annealing: short-time high-temperature annealing is performed to eliminate forging thermal stress and promote microstructure recovery; (5) Second large plastic deformation: repeat "three upsetting and three drawing", with a cumulative equivalent deformation of ≥40%, and immediately water-cooled to room temperature after forging; (6) After solution treatment at 1100-1200℃ (preferably 1150±15℃) for 0.5-1.5 h, the solution is cooled to room temperature by water, and then stabilized at 870-930℃ (preferably 920±10℃) for 1-2 h, and then cooled to room temperature in air.

[0011] Further, the other required raw materials in step (1) include: ultra-high purity iron, metallic chromium, metallic nickel, metallic manganese, silicon, metallic molybdenum, iron-niobium master alloy and iron-boron master alloy; the iron-niobium master alloy and iron-boron master alloy need to be preheated at 600-800℃ for 15-30min before smelting; the refining pressure is ≤50 Pa, and the total oxygen content of the molten steel obtained by refining is below 30 ppm; The nitrogen pressure inside the furnace for nitrogen alloying is 0.18-0.25 MPa, the molten steel temperature is 1530-1580℃, and the holding time is 10-20 min. The nitrogen protection pressure for the rare earth treatment is 0.18-0.25 MPa; the rare earth metal Ce needs to be preheated at 600-800℃ for 15-30 minutes before smelting; the temperature of the molten steel for the rare earth treatment is 1500-1550℃, and the treatment time is 3-8 minutes.

[0012] Further, the depressurization to atmospheric pressure after casting in step (2) specifically involves maintaining the nitrogen pressure in the furnace at 0.18-0.25 MPa during the initial solidification stage. After a stable solidified shell forms on the surface of the ingot and the solidification process begins, the pressure is gradually reduced to atmospheric pressure as the furnace temperature decreases, with the depressurization rate controlled at 0.01-0.03 MPa / min.

[0013] Furthermore, the single-pass deformation amount of the three upsetting and three drawing processes described in steps (3) and (5) is 15-30%, preferably 20-30%, the initial forging temperature is 1180-1250℃, and the final forging temperature is ≥1000℃.

[0014] Furthermore, in step (4), the intermediate annealing temperature is 1180-1220℃, and the holding time is 0.5-1 h / 100 mm based on the maximum effective cross-sectional size of the billet.

[0015] Furthermore, the austenitic stainless steel obtained by the above preparation method has a uniform equiaxed crystal structure and the size of the secondary precipitate is ≤1 μm; the secondary precipitate is an Nb(C,N) type MX phase.

[0016] Application of 20Cr25NiNbRE austenitic stainless steel in high-temperature components of ultra-supercritical boilers or nuclear reactors.

[0017] The design concept of this invention: This invention addresses key issues in existing 20Cr25NiNb austenitic heat-resistant steel during high-temperature service and hot working, including significant hot brittleness, continuous grain boundary precipitation and high Cr-depletion tendency at grain boundaries, and insufficient processing plasticity reserves. It proposes a novel material design method based on the synergistic effect of Ce microalloying control and large-strain microstructure reconstruction. Existing steels of this type are prone to MX phase coarsening (such as Nb(C,N)) and M under long-term high-temperature conditions. 23 C6 phase (such as Cr) 23 C6) continuously precipitates along grain boundaries, leading to Cr depletion at grain boundaries and a decrease in bonding strength, thereby inducing intergranular corrosion and hot cracking. Simultaneously, coarse inclusions formed by residual oxygen and sulfur in the molten steel significantly reduce matrix purity and high-temperature ductility. This invention, through precise control of Ce addition and its strong affinity for oxygen and sulfur, achieves deep purification and inclusion modification of the molten steel, transforming the original coarse, irregular inclusions into fine, dispersed Ce-containing rare-earth composite oxides and / or oxygen-sulfide composite inclusions, thus significantly improving material purity and microstructure uniformity. Based on this, Ce is used to regulate grain boundary energy and element diffusion behavior, altering the MX phase and M... 23 The nucleation and growth characteristics of the C6 phase inhibit the coarsening of the MX phase and promote the growth of the M phase. 23The transformation of C6 from a continuous network distribution to a dispersed state effectively blocks the grain boundary embrittlement path, reduces continuous precipitation at grain boundaries and the tendency for Cr depletion at grain boundaries, thereby improving grain boundary stability. Furthermore, by combining this with a large-strain hot deformation process, high-density dislocations and substructures are introduced through multiple passes of "three upsetting and three drawing," which not only promotes element diffusion and segregation elimination but also achieves significant grain refinement and second-phase fragmentation and redistribution, thereby improving the material's hot working safety and microstructure stability. Finally, through the synergistic mechanism of Ce microalloying purification and interface regulation, inclusion optimization, and large-strain grain refinement strengthening, the material's high-temperature ductility, hot crack resistance, and long-term service stability are further improved while maintaining high-temperature strength. This provides a new approach for designing highly reliable austenitic heat-resistant steel materials for key components of ultra-supercritical units and advanced nuclear energy systems.

[0018] Advantages and beneficial effects of the present invention: This invention provides a 20Cr25NiNbRE austenitic stainless steel and its preparation method. Its core innovation lies in the construction of an integrated design system of "ultra-high purity matrix + functionalized microalloying + microstructure reconstruction and regulation".

[0019] In terms of composition design, by employing ultra-high purity raw materials, vacuum pressurized nitrogen alloying, and rare earth deep purification technology, the content of harmful impurities such as O, S, and P is strictly controlled to achieve deep purification of the matrix, thereby suppressing stress concentration and grain boundary embrittlement induced by inclusions from the source. On this basis, an appropriate amount of Ce is introduced for microalloying (preferably 0.02-0.08 wt.%), which not only effectively refines and spheroidizes inclusions, reducing their size and harmfulness, but also optimizes grain boundary precipitation characteristics and enhances the grain boundary coordinated deformation ability by regulating grain boundary energy and element diffusion behavior through the segregation effect of grain boundaries. This significantly improves the high-temperature thermal ductility of the material and reduces the tendency of continuous precipitation and Cr depletion at grain boundaries, thus contributing to improved grain boundary stability.

[0020] In terms of preparation process, vacuum pressure melting and nitrogen control technology are used to ensure uniform composition and dense microstructure, and a multi-stage thermomechanical treatment path of "short-time annealing + two-stage upsetting and three-stage drawing" is added. High-density dislocations and substructures are introduced through large plastic deformation, which is conducive to promoting the migration of solute atoms along dislocations, subgrain boundaries and grain boundaries. At the same time, the intermediate short-time annealing realizes the release of processing thermal stress and microstructure recovery, ultimately achieving sufficient fragmentation of the cast dendritic structure and efficient homogenization of elemental segregation. Through two rounds of large strain deformation and diffusion synergy, not only are the grains significantly refined, but the Nb-rich segregation zone and related second phases are also effectively broken, improving the microstructure uniformity and stability.

[0021] Compared to a single "three-upsetting and three-drawing" hot deformation method, this invention combines intermediate short-time annealing with a second "three-upsetting and three-drawing" process. This allows the high-density dislocations, subgrain boundaries, and grain boundary diffusion channels introduced in the previous stage of deformation to recover and diffuse synergistically during the short-time high-temperature annealing. Subsequently, the second stage of large plastic deformation further breaks down the residual dendrite structure, Nb-rich segregation regions, and chain-like second phase, thereby improving the uniformity of the microstructure and the dispersion of the second phase. This multi-stage thermomechanical treatment process helps reduce local strain concentration and improves hot working stability and high-temperature ductility.

[0022] Compared to traditional processes that rely on long-term high-temperature homogenization, the "thermal deformation-diffusion synergistic homogenization" method proposed in this invention can significantly improve the thermal processing performance and subsequent service stability of materials by shortening the homogenization time, reducing energy consumption, and weakening the synergistic strengthening effect of dendrite segregation and microstructure refinement.

[0023] Compared with conventional high-temperature diffusion annealing or single-stage hot deformation processes, this invention does not rely solely on long-term high-temperature homogenization. Instead, it uses a continuous combination of the first large plastic deformation, the intermediate short-time annealing, and the second large plastic deformation to couple the dislocations, subgrain boundaries, and grain boundary diffusion channels introduced by deformation with the short-time high-temperature diffusion process. This simultaneously promotes segregation reduction, second-phase fragmentation and redistribution, and optimization of grain boundary precipitation morphology.

[0024] In summary, this invention effectively overcomes the high-temperature brittleness and structural instability of traditional 20Cr25NiNb austenitic stainless steel through the synergistic effect of component purification, Ce functionalization regulation, and innovative heat treatment processes. While maintaining high strength and oxidation resistance, it significantly improves high-temperature thermal ductility and long-term service reliability, and has significant engineering application value. Attached Figure Description

[0025] Figure 1 The image shows the EBSD image of the grain structure of the alloy steel sample from Example 1. Figure 2 This is a ternary distribution diagram of inclusion composition in the alloy steel sample of Example 1, based on ASPEX statistics; Figure 3 Cr dispersed along grain boundaries in the alloy steel sample of Example 1 23 EDS energy spectrum of C6 and corresponding Cr element; Figure 4 The image shows the Nb(C,N) type MX phase dispersed within the grains of the alloy steel sample in Example 1 and the corresponding Nb element EDS energy spectrum. Figure 5 The as-cast microstructure and EDS energy dispersive spectroscopy of the main alloying elements of the alloy steel sample in Example 2 are shown. Figure 6This is a ternary distribution diagram of inclusion composition in the alloy steel sample of Comparative Example 1, based on ASPEX statistics; Figure 7 For example, in Comparative Example 1, the Cr content in the alloy steel sample is continuously distributed along the grain boundaries. 23 EDS energy spectra of C6 and corresponding Cr elements; Figure 8 The image shows the EBSD image of the forged microstructure of the alloy steel specimen in Comparative Example 2 and its stress distribution KAM (nuclear average orientation difference) diagram. Figure 9 For the examples and comparative alloy steel samples, 10 at room temperature -3 s -1 Stress-strain curves in tensile engineering at strain rates; Figure 10 The room temperature tensile fracture morphology of the alloy steel specimens in the examples and comparative examples is shown below. Figure 10 (a) is Example 1. Figure 10 (b) is Example 2. Figure 10 (c) is Example 3. Figure 10 (d) is Example 4. Figure 10 (e) is Comparative Example 1. Figure 10 (f) is Comparative Example 2; Figure 11 For the examples and comparative examples, alloy steel samples were subjected to 10 [days / years] at 750°C. -3 s -1 Stress-strain curves in tensile engineering at strain rates; Figure 12 The examples and comparative alloy steel samples exhibit high-temperature tensile fracture morphology characteristics at 750℃; wherein, Figure 12 (a) is Example 1. Figure 12 (b) is Example 2. Figure 12 (c) is Example 3. Figure 12 (d) is Example 4. Figure 12 (e) is Comparative Example 1. Figure 12 (f) is Comparative Example 2. Detailed Implementation

[0026] The following embodiments are used to illustrate the technical solutions of the present invention and are not intended to limit the scope of protection of the present invention; those skilled in the art can make appropriate adjustments to the specific process parameters without departing from the concept of the present invention.

[0027] To make the technical solution, organizational features, and beneficial effects of the present invention clearer, specific embodiments and comparative examples are described below. Those skilled in the art should understand that appropriate adjustments can be made to the implementation methods without departing from the concept of the present invention.

[0028] Example 1 The steelmaking process is carried out using a vacuum pressure induction furnace. Raw materials are weighed according to the target composition. The raw materials selected are ultra-high purity iron, metallic chromium, metallic nickel, metallic manganese, silicon, metallic molybdenum, niobium-containing master alloy, boron-containing master alloy and rare earth metal Ce.

[0029] The preferred niobium-containing master alloy is an Fe-Nb master alloy with an Nb content of 70 wt.%; the preferred boron-containing master alloy is an Fe-B master alloy with a B content of 20 wt.%. The purity of the rare earth metal Ce is not less than 99.9 wt.%. The Fe-Nb master alloy, Fe-B master alloy, and rare earth metal Ce are crushed into small pieces of about 15 mm, coated with nickel foil, preheated at 600℃ for 20 minutes, and then taken out for use.

[0030] Before smelting, rare earth metals Ce, Fe-Nb master alloys, and Fe-B master alloys coated with preheated nickel foil are loaded into a vacuum feeding chamber. The vacuum feeding chamber is evacuated to a vacuum level of ≤50 Pa to reduce the introduction of oxygen, nitrogen, and water vapor during the feeding process.

[0031] During smelting, ultra-high purity iron and metallic nickel are first added to the crucible as the base charge. After evacuating the furnace to a pressure ≤15 Pa, power is supplied to raise the temperature. Once the base charge has melted, metallic chromium, metallic molybdenum, and Fe-Nb master alloy are added sequentially. The temperature of the molten steel is controlled at 1570±20℃ to ensure the refractory alloys are fully melted and evenly distributed. Subsequently, metallic manganese, silicon, and Fe-B master alloy are added. After the charge has completely melted, the steel is refined for 10 minutes under a vacuum condition with a pressure ≤50 Pa to reduce the gas and inclusion content in the molten steel and control the total oxygen content to below 30 ppm.

[0032] After vacuum refining, nitrogen gas with a purity of not less than 99.99% is introduced into the furnace for nitrogen alloying. The nitrogen pressure in the furnace is controlled at 0.20 MPa. The temperature of the molten steel is controlled at 1550±20℃ and held for 10 min to allow nitrogen to fully dissolve into the austenitic molten steel and improve the nitrogen yield.

[0033] After nitrogen alloying, the coated rare earth metal Ce is added to the feeding chamber through pre-vacuuming and nitrogen purging. Rare earth treatment is carried out under nitrogen pressure control at 0.20 MPa. After Ce is added, the temperature of the molten steel is controlled at 1520±20℃, held at this temperature and electromagnetically stirred for 5 minutes at a stirring speed of 20 r / min. This allows Ce to fully exert its deoxidizing, desulfurizing, and inclusion-modifying effects, transforming the inclusions in the steel into fine, dispersed Ce-containing rare earth composite oxides and / or oxysulfide composite inclusions.

[0034] During the casting stage, the nitrogen pressure inside the furnace is maintained at 0.20 MPa, and the casting temperature of the molten steel is controlled at 1520℃ to suppress nitrogen escape from the molten steel and reduce porosity and looseness defects. After casting, the nitrogen pressure inside the furnace is kept basically constant (0.18-0.25 MPa) as the temperature of the molten steel drops from 1520℃ to about 1350℃ to ensure stable nitrogen content in the molten steel and prevent nitrogen escape during the early stage of solidification. When a continuous and stable solidified shell forms on the surface of the ingot and the later stage of solidification begins, the pressure is gradually and slowly reduced to atmospheric pressure as the furnace temperature further decreases to the range of 1350-950℃. The depressurization rate is preferably controlled at 0.02 MPa / min to avoid gas evolution in the molten steel and the formation of internal defects in the ingot due to sudden pressure changes. After casting, the furnace is cooled to about 950℃ for high-temperature demolding, followed by water cooling to room temperature to obtain the ingot.

[0035] After removing the riser, ingot tail, and surface oxide scale, the ingot is shaped like a frustum with a bottom diameter of about 130 mm and a top diameter of about 160 mm. It is then homogenized at 1200℃ and held for 4.5 h.

[0036] Subsequently, the billet undergoes a first "three-upsetting and three-drawing" hot deformation process (i.e., alternating deformation processes of three upsetting and three drawing operations) at 1200℃. During this first process, the deformation amount in each pass is controlled at 18%-25%, with a cumulative equivalent deformation of approximately 65% ​​across six passes. The final forging temperature is ≥1000℃. After forging, the billet is returned to the furnace for a short-term intermediate annealing at 1200℃ (approximately 40 minutes) to eliminate thermal stress and promote microstructure recovery. A second "three-upsetting and three-drawing" hot deformation process is then performed at 1200℃. During this second process, the deformation amount in each pass is controlled at 15%-25%, with a cumulative equivalent deformation of approximately 45% across six passes. The final forging temperature is ≥1000℃. The billet is immediately water-cooled after forging to obtain a Φ60 mm forged bar.

[0037] Samples were taken longitudinally from the core of the forged bar and machined into Φ12 mm specimens. After solution treatment at 1150℃ for 45 min, the specimens were water-cooled to room temperature, followed by stabilization treatment at 930℃ for 2 h, and then cooled to room temperature in air. The measured chemical composition of the ingot is shown in Table 1.

[0038] After the above-described process, the material of Example 1 has the following significant structural characteristics: (1) The grain structure is uniform equiaxed crystal with an average grain size of 26.7 μm. There is no obvious mixed coarse and fine crystal phenomenon, as follows: Figure 1 ; (2) The composition of inclusions in steel was detected using the ASPEX automatic inclusion analysis system, and the statistical analysis results are as follows: Figure 2 As shown, the inclusions are mainly fine, dispersed cerium oxides with an average size of 2 μm; (3) M at the grain boundary 23The C6 phase exhibits a discontinuous and discrete distribution, failing to form a continuous network structure, as follows: Figure 3 As shown; (4) The Nb(C,N) type MX phase is diffusely distributed and has a small size, significantly smaller than 1 μm, as shown below. Figure 4 As shown; The above results indicate that Ce plays a synergistic role in inclusion modification and grain boundary regulation in steel.

[0039] Example 2 The process of Example 2 is the same as that of Example 1, except that the control range of Ce content is different (see Table 1 for specific composition). In order to obtain an initial grain structure with similar size, the Φ12 mm sample was solution treated at 1150℃ for 47 min, and then subjected to the same stabilization treatment. The average grain size of the material obtained after the above process is 25.5 μm.

[0040] Example 3 The process of Example 3 is the same as that of Example 1, except that the control range of Ce content is different (see Table 1 for specific composition). In order to obtain an initial grain structure with similar size, the Φ12 mm sample was solution treated at 1150℃ for 52 min, and then the same stabilization treatment was performed. The average grain size of the material obtained after the above process is 25.9 μm.

[0041] Example 4 The process of Example 4 is the same as that of Example 1, except that the control range of Ce content is different (see Table 1 for specific composition). In order to obtain an initial grain structure with similar size, the Φ12 mm sample was solution treated at 1150℃ for 72 min, and then subjected to the same stabilization treatment. The average grain size of the material obtained after the above process is 25.1 μm.

[0042] EDS analysis of the samples in Examples 2-4 showed that, under Ce-containing conditions, no obvious macroscopic segregation was observed in the as-cast microstructure, and the inclusions were finely and dispersedly distributed. The EDS energy dispersive spectroscopy results of the main alloying elements in the as-cast microstructure of Example 2 are as follows: Figure 5 As shown.

[0043] Comparative Example 1 Comparative Example 1 and Example 1 used the same preparation process, except that there was no rare earth treatment process, that is, Ce was not added during the steel smelting process. The specific composition is shown in Table 1. After the above process, the average grain size of the material was 26.5 μm.

[0044] Its organizational characteristics are as follows: (1) The inclusions are unevenly distributed in size, and the main type is hard alumina inclusions. The ASPEX statistical analysis results are as follows: Figure 6 As shown; (2) Locally continuous distribution of M exists at the grain boundaries 23 C6 phase, such as Figure 7 As shown; (3) The uniformity of the microstructure in the grain boundary region is poor; The above results indicate that when Ce is lacking in the steel, the grain boundary control ability is significantly reduced.

[0045] Comparative Example 2 Comparative Example 2 used commercially available 20Cr25NiNb austenitic stainless steel. This material was prepared using a conventional high-temperature diffusion annealing combined with triaxial forging process. Specifically, the ingot was heated to 1200℃ and held for more than 5 hours to promote the diffusion and homogenization of as-cast dendrite segregation and primary precipitates. Subsequently, triaxial forging was performed in the high-temperature austenitic region. During the forging process, the billet was deformed in multiple passes along different directions, with a cumulative forging ratio greater than 7, and the billet temperature during the forging process was not lower than 950℃. After forging, the material was immediately water-cooled to room temperature to obtain the forged material.

[0046] It should be noted that although Comparative Example 2 underwent high-temperature diffusion annealing and large forging ratio triaxial forging treatment, it did not adopt the two-stage high plasticity hot working and intermediate annealing synergistic process described in this invention, which is "homogenization treatment - first three upsetting and three drawing - intermediate annealing - second three upsetting and three drawing".

[0047] To eliminate the influence of grain size differences on performance, a Φ12 mm sample of the material was subjected to solution treatment at 1200℃ for 2 h and stabilization treatment at 930℃ for 2 h. The average grain size of the material was found to be 28.9 μm, which is basically equivalent to the average grain size of the materials in Examples 1-4. The measured composition is shown in Table 1.

[0048] EBSD and KAM analysis results are as follows Figure 8 As shown, the microstructure of Comparative Example 2 still exhibits localized uneven deformation, uneven residual strain distribution, and chain-like MX phase distribution characteristics. This indicates that when only high-temperature diffusion annealing and triaxial forging processes are used, the as-cast / forged genetic microstructure, Nb-rich segregation zone, and chain-like distribution of the second phase cannot be sufficiently improved. However, the present invention, through the synergistic effect of two-stage "three upsetting and three drawing" and intermediate short-time annealing, can more effectively promote microstructure homogenization and the dispersed distribution of the second phase.

[0049] Comparative Example 2, using commercially available 20Cr25NiNb austenitic stainless steel, differs in composition from the materials in the examples. It primarily serves to illustrate that under conventional high-temperature diffusion annealing combined with triaxial forging, similar materials may still exhibit problems such as chain-like MX phases, uneven residual strain, and insufficient high-temperature plasticity. Comparative Example 1, on the other hand, compares the effect of Ce addition on microstructure and properties under similar composition and hot-working conditions. Figure 8 It is known that using only conventional high-temperature diffusion annealing and triaxial forging processes is still insufficient to fully improve the distribution of Nb-rich segregation zones and chain-like second phases. The two-stage "three upsetting and three drawing" combined with intermediate short-time annealing process of this invention is more conducive to promoting the homogenization of the microstructure and the diffuse distribution of the second phase.

[0050] Table 1. Chemical composition of austenitic stainless steel materials in the examples and comparative examples, with the balance being Fe (wt.%).

[0051] Performance testing and results analysis Under conditions of essentially uniform grain size (approximately 27±4 μm), tensile properties were tested on the specimens of the examples and comparative examples. The room temperature tensile testing standard was GB / T 228.1-2021 "Metallic materials, tensile testing—Part 1: Test at room temperature," while the high temperature tensile testing standard was GB / T 228.2-2015 "Metallic materials, tensile testing—Part 2: Test at high temperature." The strain rate for both tests was 1×10⁻⁶. -3 s -1 The results are as follows: (1) Room temperature performance ( Figure 9 ) The yield strengths of Examples 1-4 are approximately 362 MPa, 352 MPa, 349 MPa, and 346 MPa, respectively, while those of Comparative Examples 1 and 2 are approximately 355 MPa and 325 MPa, respectively. Except for Comparative Example 2, which did not employ the hot-working process of this invention, the yield strengths of the examples are similar to those of Comparative Example 1, indicating that under conditions of similar grain size, this invention does not significantly reduce the yield strength level of the material.

[0052] The tensile strengths of Examples 1-4 were approximately 738 MPa, 723 MPa, 714 MPa, and 706 MPa, respectively, while those of Comparative Examples 1 and 2 were approximately 734 MPa and 674 MPa, respectively. The results indicate that the embodiments of the present invention, while maintaining a high strength level, exhibit better overall tensile performance than the material of Comparative Example 2, which did not employ the microstructure control process of the present invention.

[0053] Furthermore, the elongation rates of Examples 1-4 reached 43%, 45%, 46%, and 46%, respectively, and the reduction of area reached 58%, 52%, 59%, and 61%, respectively; in contrast, the elongation rates of Comparative Examples 1 and 2 were 43% and 38%, respectively, and the reduction of area were only 46% and 44%, respectively. The results show that the embodiments of the present invention significantly improve the plastic deformation capacity and the energy dissipation capacity before fracture while maintaining the basic stability of strength.

[0054] At the same time, by Figure 10 As can be seen, the fracture surfaces of Examples 1-4 exhibit more pronounced and uniformly distributed dimples and shear lips, indicating that the materials possess superior ductile fracture behavior; while the materials of Comparative Examples 1-2 show a more obvious tendency towards brittle fracture. These results demonstrate that the present invention effectively improves the material's room temperature deformation compatibility and damage resistance through microstructure regulation.

[0055] (2) High temperature performance at 750℃ Figure 11 ) At 750°C, the yield strengths of Examples 1-4 were approximately 184 MPa, 199 MPa, 189 MPa and 175 MPa, respectively, while those of Comparative Example 1 and Comparative Example 2 were approximately 186 MPa and 187 MPa, respectively. The differences between the samples were small, indicating that the present invention did not have a significant adverse effect on the high-temperature load-bearing capacity of the material.

[0056] The tensile strengths of Examples 1-4 reached 365 MPa, 354 MPa, 353 MPa, and 345 MPa, respectively, while those of Comparative Examples 1 and 2 were 356 MPa and 307 MPa, respectively. The results show that the embodiments of the present invention maintain high tensile properties under high-temperature conditions, and in particular, exhibit superior high-temperature strength retention compared to the material of Comparative Example 2, which did not employ the hot-working path of the present invention.

[0057] Furthermore, the elongation rates of Examples 1-4 reached 53%, 62%, 66%, and 59%, respectively, and the reduction of area reached 81%, 85%, 86%, and 82%, respectively; while the elongation rates of Comparative Examples 1 and 2 were 56% and 62%, respectively, and the reduction of area were only 71% and 55%, respectively. It is evident that the embodiments of the present invention exhibit superior high-temperature plasticity and thermal ductility, especially with a significantly improved reduction of area, indicating that the material possesses stronger strain compatibility and resistance to thermal cracking during high-temperature deformation.

[0058] Combination Figure 12Fracture morphology analysis shows that the materials of Examples 1-4 exhibit more complete dimple aggregation and plastic tearing characteristics after high-temperature tensile testing. The fracture surface of Comparative Example 1 is relatively flat, with a significant reduction in the number of dimples. Quasi-cleavage features and uneven plastic deformation can be observed in local areas. In contrast, the fracture surface of Comparative Example 2 shows more obvious coarse pores and loose tearing characteristics. The fracture structure is rough and the plastic deformation is uneven, indicating that the material is more prone to local strain concentration and rapid crack propagation during high-temperature tensile testing. This indicates that Examples 1-4 have superior high-temperature toughness fracture behavior.

[0059] In summary, the material of this invention achieves a significant improvement in high-temperature plasticity and thermal ductility while maintaining relatively stable high-temperature strength.

[0060] Mechanism Explanation Under conditions of essentially uniform grain size, the yield strength difference between the examples and the comparative examples is small, indicating that the contribution of fine grain strengthening is basically equivalent. Therefore, the performance improvement of this invention mainly comes from improved microstructure stability and regulation of grain boundary behavior.

[0061] The performance improvement mechanism of this invention can be summarized as follows: (1) Optimization effect of inclusions Ce forms fine, dispersed Ce-containing rare earth composite oxides and / or oxysulfides with oxygen and sulfur elements, which reduces the size of the inclusions and makes them more uniformly distributed, thereby reducing local stress concentration and crack initiation tendency.

[0062] (2) Regulation of grain boundary precipitation Ce modulates grain boundary energy and element diffusion behavior, enabling M 23 The C6 precipitate changes from a continuous network distribution to a discontinuous discrete distribution, thereby reducing grain boundary embrittlement and Cr depletion effects and improving grain boundary stability.

[0063] (3) Precipitation enhancement effect Fine and dispersed Nb(C,N) type MX phases increase the resistance to dislocation movement and suppress damage accumulation during high-temperature deformation, thereby improving the material's strength and toughness synergy.

[0064] (4) Synergistic effect of two-stage thermomechanical treatment The two-stage “three-upsetting and three-drawing” combined with intermediate short-time annealing can further break up dendritic segregation regions and Nb-rich regions, promote uniform diffusion of elements and uniform distribution of the second phase, reduce local residual strain concentration, and improve the uniformity of the microstructure and the ability to coordinate high-temperature deformation.

[0065] The synergistic effect of the above factors enables the material to simultaneously improve tensile strength, high-temperature thermal ductility, and damage resistance while maintaining a basically stable yield strength.

[0066] In summary, by controlling Ce microalloying and employing a synergistic process of "two-stage large plastic hot deformation + intermediate annealing," this invention achieves inclusion refinement, grain boundary precipitation optimization, and microstructure homogenization control. This results in 20Cr25NiNbRE austenitic stainless steel that significantly improves high-temperature thermal ductility and damage resistance while maintaining high-temperature strength, thereby enhancing the material's service reliability in ultra-supercritical boilers and high-temperature components of advanced nuclear energy systems.

[0067] This material has promising applications in ultra-supercritical boilers and high-temperature components for nuclear power plants.

Claims

1. A method for preparing 20Cr25NiNbRE austenitic stainless steel, characterized in that, Includes the following steps: (1) Smelting, nitrogen alloying and rare earth treatment: Add the required raw materials other than rare earth metal Ce to the vacuum pressure induction furnace according to the target composition. After refining at 1550-1600℃ for 10-15 minutes, nitrogen gas is introduced into the furnace for nitrogen alloying. Then, rare earth metal Ce is added under nitrogen protection for rare earth treatment to obtain molten steel to be poured. (2) Casting and homogenization treatment: molten steel is cast at 1500-1540℃ with nitrogen pressure maintained at 0.18-0.25 MPa; after casting, the pressure is released to atmospheric pressure, the furnace is cooled to 900-950℃ and then demolded at high temperature, followed by water cooling to room temperature to obtain ingots; the ingots are heated to 1180-1250℃ for homogenization treatment and held for 4-6 h. (3) First large plastic deformation: The ingot is subjected to the first "three upsetting and three drawing" in the single-phase austenite temperature range of 1180-1250℃, and the cumulative equivalent deformation is ≥60%; (4) Intermediate annealing: short-time high-temperature annealing is performed to eliminate forging thermal stress and promote microstructure recovery; the intermediate annealing temperature is 1180-1220℃, and the holding time is 0.5-1 h / 100 mm based on the maximum effective cross-sectional size of the billet. (5) Second large plastic deformation: repeat "three upsetting and three drawing", with a cumulative equivalent deformation of ≥40%, and immediately water-cooled to room temperature after forging; (6) After solution treatment at 1100-1200℃ for 0.5-1.5 h, water-cool to room temperature, then stabilize at 870-930℃ for 1-2 h, and then cool to room temperature in air; The target component, by mass percentage, has the following chemical composition: C ≤ 0.10%, Cr: 19-21%, Ni: 24-26%, Mn: 0.3-1.5%, Si: 0.1-0.5%, Mo: 1-2%, Nb ≤ 0.30%, N: 0.10-0.30%, B: 0.001-0.003%, Ce: 0.005-0.12%, Al ≤ 0.05%, Ti ≤ 0.02%, P ≤ 0.03%, S ≤ 0.02%, O ≤ 0.003%, H ≤ 0.0002%, with the balance being Fe and unavoidable impurities; The deformation per pass of the three upsetting and three drawing processes described in steps (3) and (5) is 15-30%, the initial forging temperature is 1180-1250℃, and the final forging temperature is ≥1000℃.

2. The method for preparing 20Cr25NiNbRE austenitic stainless steel according to claim 1, characterized in that: Other required raw materials in step (1) include: ultra-high purity iron, metallic chromium, metallic nickel, metallic manganese, silicon, metallic molybdenum, iron-niobium master alloy and iron-boron master alloy; the iron-niobium master alloy and iron-boron master alloy need to be preheated at 600-800℃ for 15-30min before smelting; the refining pressure is ≤50 Pa, and the total oxygen content of the molten steel obtained from refining is below 30 ppm; The nitrogen pressure inside the furnace for nitrogen alloying is 0.18-0.25 MPa, the molten steel temperature is 1530-1580℃, and the holding time is 10-20 min. The nitrogen protection pressure for the rare earth treatment is 0.18-0.25 MPa; the rare earth metal Ce needs to be preheated at 600-800℃ for 15-30 minutes before smelting; the temperature of the molten steel for the rare earth treatment is 1500-1550℃, and the treatment time is 3-8 minutes. The specific steps of depressurizing to atmospheric pressure after casting in step (2) are as follows: In the early stage of solidification, the nitrogen pressure in the furnace is maintained at 0.18-0.25 MPa. After a stable solidified shell forms on the surface of the ingot and the solidification process begins, the pressure is gradually reduced to atmospheric pressure as the furnace temperature decreases. The depressurization rate is controlled at 0.01-0.03 MPa / min.

3. The method for preparing 20Cr25NiNbRE austenitic stainless steel according to claim 1, characterized in that: The austenitic stainless steel obtained by this preparation method has a uniform equiaxed crystal structure and the size of the secondary precipitates is ≤1 μm; the secondary precipitates are Nb(C,N) type MX phases.

4. A 20Cr25NiNbRE austenitic stainless steel, characterized in that: The austenitic stainless steel is prepared according to the preparation method described in any one of claims 1-3; the chemical composition of the stainless steel by mass percentage is: C ≤ 0.10%, Cr: 19-21%, Ni: 24-26%, Mn: 0.3-1.5%, Si: 0.1-0.5%, Mo: 1-2%, Nb ≤ 0.30%, N: 0.10-0.30%, B: 0.001-0.003%, Ce: 0.005-0.12%, Al ≤ 0.05%, Ti ≤ 0.02%, P ≤ 0.03%, S ≤ 0.02%, O ≤ 0.003%, H ≤ 0.0002%, with the balance being Fe and unavoidable impurities.

5. The 20Cr25NiNbRE austenitic stainless steel according to claim 4, characterized in that: The chemical composition, by mass percentage, further satisfies the following: Nb: 0.15-0.25%, N: 0.15-0.25%, C: 0.03-0.07%, Ti ≤ 0.01%.

6. The 20Cr25NiNbRE austenitic stainless steel according to claim 4, characterized in that: The chemical composition, by mass percentage, further satisfies the following: P ≤ 0.01%, S ≤ 0.005%, O ≤ 0.002%, H ≤ 0.0002%.

7. The 20Cr25NiNbRE austenitic stainless steel according to claim 4, characterized in that: The chemical composition, by mass percentage, further satisfies: Ce: 0.02-0.08%.

8. The application of the 20Cr25NiNbRE austenitic stainless steel according to any one of claims 4-7 in high-temperature components of ultra-supercritical boilers or nuclear reactors.

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