A high-temperature-resistant low-activation steel and a preparation method thereof

Abstract

This invention relates to the field of metallic materials and hot working technology. The invention discloses a high-temperature resistant, low-activation steel and its preparation method. The chemical composition of the raw materials is as follows (by mass percentage): Cr: 8.0–10.0%, W: 1.0–2.0%, V: 0.2–0.5%, Ta: 0.05–0.2%, Y: 0.005–0.02%, Ce: 0.001–0.005%, B: 0.001–0.005%, C: 0.08–0.16%, N: 0.015–0.035%, Mn: 0.3–0.8%, with the balance being Fe. The steel matrix is ​​tempered martensite, with the original austenite grains having an equivalent size of less than 10 μm, and containing dispersed MX-type carbides (nitrides) and M... 23 C6 type carbides. The preparation process involves vacuum induction melting, homogenization and forging, multi-cycle thermomechanical treatment (TMT), and tempering. This significantly refines the austenite grains and martensite substructure, thereby enhancing the synergistic effect of strength and toughness. The material achieves excellent tensile strength at 650 °C while maintaining an impact absorption energy greater than 150 J at room temperature and -20 °C, making it suitable for use in nuclear fusion reactor core structural materials.

Description

Technical Field

[0001] This invention relates to the field of metal materials and heat treatment technology, specifically to a high-temperature resistant, low-activation steel and its preparation method, and more specifically to a thermomechanical treatment preparation method for a high-temperature resistant, low-activation ferrite / martensitic (RAFM) steel suitable for tritium production blanket structures in fusion reactors. Background Technology

[0002] The statements in this section provide only background information relevant to the disclosure of this application and may not constitute prior art.

[0003] With the global energy structure undergoing a green and low-carbon transformation, nuclear energy, especially nuclear fusion energy, has become a promising next-generation clean energy source. my country is accelerating its research and development of clean energy technologies, including fusion energy. Currently, as nuclear fusion energy is gradually shifting from combustion experiments to the construction of engineering experimental reactors, the performance requirements for core structural materials are far higher than those of existing nuclear fission reactors.

[0004] Low-activation ferritic / martensitic steel (RAFM steel) is the preferred material for tritium breeder blankets in fusion reactors due to its superior thermal conductivity, low expansion, and resistance to radiation swelling. However, the service temperature of current mainstream RAFM steels (such as EUROFER97, F82H, and CLF-1) is still limited to a maximum of 550 °C, which is insufficient to meet the high-temperature structural material requirements of future fusion reactors' high-heat-transfer-efficiency breeder blankets, such as low-activation steels resistant to 650 °C. Furthermore, existing RAFM steels still face bottlenecks in mechanical stability, radiation resistance, and weldability under neutron irradiation, hot corrosion, and high-temperature creep environments. In particular, the strength loss and toughness degradation caused by decreased microstructure stability in the weld heat-affected zone and under high-temperature service conditions urgently need to be addressed through novel process routes and alloy design methods.

[0005] Recent studies have shown that by constructing multi-scale microstructures during thermomechanical treatment (TMT) to induce the precipitation of MX-type nanophases, and combining this with a multi-cycle quench-deformation-tempering process, it is possible to achieve refined and stable alloy microstructures and significantly improve the high-temperature strength and low-temperature toughness of the material. However, how to achieve engineerable and large-scale preparation under the traditional smelting-rolling system remains a key technical challenge that needs to be overcome in the development of RAFM steel. Summary of the Invention

[0006] The purpose of this invention is to address the limitations of current nuclear fusion reactor core materials in terms of high-temperature resistance, as well as strength loss and toughness degradation during service. This invention provides a high-temperature resistant, low-activation steel and its preparation method. A method for preparing RAFM steel with both high-temperature tensile properties above 650 °C and room-temperature / low-temperature impact toughness is proposed. By combining a multi-cycle TMT process with a precipitation control strategy, the contradiction between performance and processing of fusion reactor blanket structural materials is resolved, filling a gap in existing technology. The obtained material exhibits good tensile properties at 650 °C and impact energy ≥150 J at room temperature and -20 °C.

[0007] The technical solution of the present invention is as follows: This invention provides a high-temperature resistant, low-activation steel. The chemical composition of the steel, by mass percentage, is: Cr: 8.0–10.0%, W: 1.0–2.0%, V: 0.2–0.5%, Ta: 0.05–0.2%, Y: 0.005–0.02%, Ce: 0.001–0.005%, B: 0.001–0.005%, C: 0.08–0.16%, N: 0.015–0.035%, Mn: 0.3–0.8%, with Fe as the balance. The steel's matrix is ​​tempered martensite, with an equivalent original austenite grain size of less than 10 μm, and contains dispersed MX-type carbides (nitrides) and M... 23 C6 type carbides.

[0008] According to a preferred embodiment, the mass fraction of highly activating elements and impurity elements in the steel is limited as follows: Ni≤0.02%, Co≤0.01%, Nb≤0.01%, Mo≤0.01%, Cu≤0.01%, Al≤0.05%, Si≤0.10%, P≤0.005%, S≤0.005%, O≤0.005%.

[0009] According to a preferred embodiment, the MX-type carbide (nitride) is a carbide (nitride) mainly composed of Ta and V, with an average size of 10–60 nm; the M 23 C6 carbides have an average size of 50–200 nm, and are diffusely distributed along the grain interior and grain boundaries, forming a discontinuous network.

[0010] According to a preferred embodiment, the tensile strength of the steel material at 650 °C R m ≥400 MPa and yield strength (or specified plastic elongation strength) R p0.2 ≥350 MPa, elongation after fracture A≥20%; impact absorption energy at room temperature and / or -20 ℃ A KV2 ≥150 J.

[0011] More preferably, the yield strength of the steel material at 650 °C R p0.2 ≥400 MPa, tensile strength R m ≥460MPa; at room temperature and / or Impact absorption energy at 20℃ A KV2 The value is 150–250 J.

[0012] Another aspect of the present invention provides the application of a high-temperature resistant, low-activation steel as a structural material for nuclear fusion reactor cores, as described above.

[0013] Preferably, for example, applications in structures such as the first wall / tritium-producing blanket of a fusion reactor, divertors, and their pressure-bearing / heat-conducting components.

[0014] Preferably, the steel is supplied in one or more of the following forms: plate, bar, or strip.

[0015] Another aspect of the present invention provides a method for preparing a high-temperature resistant, low-activation steel as described above, comprising the following steps: (1) Vacuum induction melting: Melting is carried out under vacuum degree ≤1 Pa. The elements are added in order of melting point from high to low. Trace elements such as Y, Ce and B are added in the later stage by lifting the material. The components are uniformly stirred by electromagnetic stirring and then cast into steel ingots. The duration of electromagnetic stirring is 0.5 to 1.5 h, preferably about 1 h.

[0016] (2) Homogenization and forging: After the steel ingot is homogenized by holding at 1100~1250 ℃ for ≥5 h, it is forged by alternating upsetting and drawing at 1100~1200 ℃ for ≥3 times, with a forging ratio of ≥4 and a final forging temperature of ≥850 ℃; (3) Multi-cycle thermomechanical treatment (TMT): First round: Heat the forged billet after step (2) to 1000-1150 ℃ and hold for ≥30 min, then perform hot deformation at 750-800 ℃ with a deformation amount ≤30%, and then cool to room temperature by water or spray; Circulation round: Heat the billet again to 900-1000 ℃ and hold for ≥20 min, then perform hot deformation at 750-800 ℃ with a deformation amount of 35-60%, and then cool to room temperature by water or spray; The circulation round is performed at least once, preferably two to three times; The above process can further refine the original austenite grains and promote the formation of fine dispersed MX phase, especially enhance the precipitation effect of nano carbonitrides formed by elements such as Ta and V.

[0017] (4) Tempering: After holding at 700-760 ℃ for 30-90 min, air cool or air cool to room temperature; finally, tempered martensite matrix, fine original austenite grains and dispersed MX and M are obtained. 23 High-temperature, low-activation steel with C6 precipitates. This tempering process helps stabilize the martensitic structure, release thermal stress, and promote the secondary precipitation and pinning effect of MX-type precipitates, thereby significantly improving high-temperature creep strength and room-temperature impact toughness.

[0018] According to a preferred embodiment, the hot deformation of each wheel in the TMT process is carried out by hot compression or hot rolling, and after the hot deformation of each wheel is completed, it is directly and rapidly cooled to room temperature without isothermal holding.

[0019] According to a preferred embodiment, the holding time for tempering in step (4) is 45 to 90 minutes.

[0020] According to a preferred embodiment, the steel, after tempering, acquires a stable lath martensite structure with fine subgrain / dislocation cell structure, and the precipitates have a pinning effect on grain boundary migration and dislocation movement. After two or more rounds of TMT and tempering, a microstructure with an average original austenite grain equivalent size of 3-8 μm is obtained. By introducing high dislocation density and high-energy interfaces during the TMT process, the preferential growth of MX nanoprecipitates formed by V and Ta is promoted within the grains, and they interact with the M precipitates formed during tempering. 23 C6 synergistically achieves pinning of grain boundaries and dislocations. The yield strength of the steel prepared by the described method at 650 °C is... R p0.2 ≥400 MPa, tensile strength R m ≥460 MPa; at room temperature and / or Impact absorption energy at 20℃ A KV2 The value is 150–250 J.

[0021] This invention optimizes the alloy composition system by introducing strong MX-type carbonitride forming elements (such as Ta and V) to promote the dispersed precipitation of nanoscale MX phases during TMT thermomechanical treatment, thereby enhancing high-temperature strength. Through compositional control, the austenitization transformation temperature and Ms point of the alloy are controlled, ensuring a stable tempered martensite structure after quenching and effectively suppressing the residual ferrite structure at high temperatures. Combined with a multi-pass TMT process and a cyclic austenitization + deformation + quenching path, significant refinement of austenite grains and martensite substructure is achieved, thus enhancing the synergistic effect of strength and toughness. Finally, tempering at a reasonable temperature (700~760 °C) stabilizes and redistributes the nanoscale precipitates in the tempered martensite, enabling the material to achieve excellent tensile strength at 650 °C while maintaining an impact absorption energy greater than 150 J at room temperature and -20 °C, meeting the operational requirements of fusion reactors.

[0022] Specifically, the alloy first undergoes high-temperature homogenization and multi-pass forging to obtain a uniform matrix structure. Subsequently, it undergoes two or more austenitizing-hot deformation cycles (TMT process) to introduce a large number of dislocations and intragranular nucleation regions at different strain levels, while further controlling the austenite grain size to stabilize it at a sub-fine grain level. During this process, the deformation-induced dislocation-dense regions and high-energy interfaces significantly promote the simultaneous precipitation of MX phases such as Ta and V during deformation. These precipitates not only pin dislocations and stabilize the substructure but also form a dispersed distribution through heterogeneous nucleation at grain boundaries and within grains, resulting in a high coupling between the strengthening mechanism and the microstructure evolution process. Finally, tempering at 700–760 °C further stabilizes the microstructure and optimizes the MX precipitate state and distribution, achieving a synergistic effect of precipitation strengthening and substructure refinement. The resulting microstructure is tempered martensite with uniform, fine proto-austenite grains and densely distributed MX nanophases. The material exhibits excellent high-temperature creep strength, good plasticity, good radiation stability, and resistance to helium bubble formation.

[0023] Compared with existing technologies, the advantages of this invention are: 1. A high-temperature resistant, low-activation steel, which, by introducing appropriate amounts of Cr, W, Ta, V, Y, and B in its composition design, synergistically optimizes the alloy's thermal stability, carbonitride precipitation behavior, and martensitic microstructure stability. Cr and W enhance high-temperature strength and oxidation resistance; Ta and V provide precursors for the formation of stable nanoscale MX-type precipitates; and Y and B help refine grains and suppress early fracture induced by inclusions. Based on this, by optimizing and expanding the traditional TMT thermomechanical treatment process, a multi-pass austenitization + hot deformation cycle is constructed to achieve synergistic refinement of the original austenite grains and martensitic substructure. 2. A method for preparing high-temperature resistant, low-activation steel involves subjecting a homogenized forging billet to multiple TMT processes and cyclic austenitization processes, combined with tempering at 700~760 °C. Through the synergistic design strategy of composition, microstructure, and process, the resulting low-activation martensitic steel exhibits significantly superior comprehensive performance. This method also expands the application boundaries of TMT processes in the development of high-temperature structural materials, providing a new material solution for next-generation advanced nuclear energy systems.

[0024] 3. The application of a high-temperature resistant, low-activation steel, and the resulting yield strength of the low-activation martensitic steel. R p0.2 ≥400 MPa, tensile strength R m ≥460 MPa; at room temperature and / or Impact absorption energy at 20℃ A KV2 The J value is 150–250; it meets multiple stringent requirements for fusion reactor blanket structure materials, such as low activation, high temperature resistance, radiation resistance, and good toughness, and can be applied to structures such as the first wall / tritium-producing blanket, divertor, and their pressure-bearing / heat-conducting components in nuclear fusion reactors. Attached Figure Description

[0025] Figure 1 This is a process flow diagram for a high-temperature resistant, low-activation steel. Figure 2 The room temperature (25 °C) tensile curves of samples from Examples 1-4 are shown. Figure 3 High-temperature (650 °C) tensile curves of samples from Examples 1-4; Figure 4-7 The optical electron microscopy (OEM) morphology of the high-temperature resistant, low-activation steel samples prepared in Examples 1-4 are shown respectively; the scale bars in the figures are 50 μm (left) and 20 μm (right). Figure 8 The SEM and TEM microstructures of the high-temperature resistant, low-activation steel sample prepared in Example 4 are shown. Detailed Implementation

[0026] The specific embodiments listed in this invention are merely examples, and the invention is not limited to the specific embodiments described below. For those skilled in the art, any equivalent modifications and substitutions to the embodiments described below are also within the scope of this invention. Therefore, all equivalent transformations and modifications made without departing from the spirit and scope of this invention should be covered within its scope. Where specific conditions are not specified in the embodiments, conventional conditions or conditions recommended by the manufacturer shall apply. All reagents or instruments whose manufacturers are not specified are commercially available conventional products. To better illustrate this invention, numerous specific details are provided in the following detailed embodiments. Those skilled in the art should understand that this invention can be practiced even without certain specific details. In other embodiments, methods, means, equipment, and steps well known to those skilled in the art are not described in detail in order to highlight the main points of this invention.

[0027] Unless otherwise defined, all technical and scientific terms used herein have the same meanings as commonly understood by one of ordinary skill in the art. Unless otherwise specified, all units used in this specification are International Standard Units (SI), and all numerical values ​​and ranges appearing in this invention should be understood to include systematic errors unavoidable in industrial production.

[0028] The features and performance of the present invention will be further described in detail below with reference to embodiments.

[0029] Example 1 (Comparative example, without TMT or NT processes) The nominal composition of the alloy, by mass percentage, is: Cr 8.5%, W 1.5%, V 0.35%, Ta 0.12%, Y 0.01%, B 0.003%, Ce 0.003%, C 0.12%, N 0.025%, Mn 0.56%, Fe balance.

[0030] refer to Figure 1 The preparation process is carried out according to the following steps: (1) Batching and Smelting: Based on the above composition, the raw materials were converted into mass percentages, and all alloying elements were carefully weighed and their surfaces cleaned. Vacuum induction melting was adopted, maintaining a vacuum degree ≤1 Pa in the furnace. Elements were added to the crucible in descending order of melting point, with trace elements such as Y, Ce, and B added later by a hoisting method. After smelting, electromagnetic stirring was performed for 1 hour to ensure uniform alloy composition. After smelting, the alloy was cast into steel ingots of Φ 300 mm × 500 mm.

[0031] (2) Homogenization and forging: After the ingot is naturally cooled to room temperature, it is homogenized by holding it at 1200 ℃ for 5 hours to eliminate chemical segregation. After homogenization, it is forged. The initial forging temperature is about 1150 ℃. It is subjected to three repeated "upsetting-drawing" hot working processes with a forging ratio of 4. The final forging temperature is about 850 ℃. After forging, it is air-cooled to room temperature.

[0032] (3) Normalizing and tempering: After forging, the sample was subjected to normalizing heat treatment: 960 ℃ for 40 minutes, water cooling; then tempering treatment was performed, 740 ℃ for 60 minutes, air cooling.

[0033] Example 1 is a comparative example without thermomechanical treatment; its final microstructure was obtained through forging, normalizing, and tempering processes. The tempered microstructure is as follows: Figure 4 As shown, the material exhibits a relatively large overall grain size, with an equivalent grain size of approximately 50 μm, clear grain boundary contours, and a limited number of substructures within the grains. In this embodiment, the matrix microstructure is dominated by tempered martensite with large lath dimensions, relatively low dislocation density, and a limited number of carbides formed during tempering, which are relatively coarse and dispersed. These microstructural characteristics indicate that, without the introduction of thermomechanical deformation refinement and dislocation control, conventional heat treatment alone is insufficient to obtain a high-density substructure and a stable precipitation-strengthened structure in this alloy system. This microstructure directly results in lower strength levels at both room and high temperatures in this embodiment; however, due to good grain boundary continuity and low internal stress concentration, its impact toughness remains at a high level.

[0034] Example 2 The alloy composition is the same as in Example 1.

[0035] (1) Batching and smelting: Same as in Example 1.

[0036] (2) Homogenization and forging: Same as in Example 1.

[0037] (3) Single-round thermomechanical treatment (TMT1): After forging, the sample is heated to 1050 ℃ and held for 30 min. Then, it is hot-compressed or rolled at 760 ℃, and the deformation is controlled at 40%. After the hot deformation is completed, it is immediately cooled to room temperature by water cooling or spraying.

[0038] (4) Tempering treatment: The TMT-treated sample was tempered at 740 °C for 60 min and then air-cooled.

[0039] Example 2 introduces a thermomechanical treatment based on forging. The tempered microstructure is as follows: Figure 5As shown, compared to Example 1, the grain size is significantly reduced, and the grains are elongated to a certain extent along the hot deformation direction, resulting in a significantly refined equivalent grain size. The hot deformation introduced during a single TMT process causes a large number of dislocations and subgrain structures to form inside the grains. Some of these substructures are still retained after tempering, thus significantly improving the yield strength and tensile strength of the material. However, due to only one TMT process, the microstructure still exhibits strong orientation and inhomogeneity, with high dislocation density and significant stress concentration in local areas. Simultaneously, the amount of carbides precipitated during tempering is limited, resulting in insufficient interface control. While this microstructure improves strength, it also negatively impacts toughness, leading to a decrease in impact performance compared to Example 1.

[0040] Example 3 The alloy composition is the same as in Example 1.

[0041] (1) Batching and smelting: Same as in Example 1.

[0042] (2) Homogenization and forging: Same as in Example 1.

[0043] (3) Multi-cycle thermomechanical treatment (TMT2): The forged specimen is first heated to 1050 ℃ and held for 30 min. Then, the first round of hot deformation is carried out at 780 ℃, using hot compression or rolling process, with a deformation amount of 20%. After hot deformation, the specimen is immediately cooled to room temperature by water cooling or spray cooling.

[0044] After cooling to room temperature, the sample was heated again to 950 ℃ and held for 20 minutes. Then, a second round of hot deformation was carried out at 760 ℃ with a deformation amount of 50%. After the deformation was completed, the sample was immediately water-cooled to room temperature.

[0045] (4) Tempering treatment: After TMT treatment, the sample was tempered at 740 °C for 60 min and then air-cooled.

[0046] Example 3 further modulates the microstructure through two rounds of cyclic thermomechanical treatment, resulting in a tempered microstructure as follows: Figure 6 As shown in the figure, compared with Example 2, the grain size is further refined, the grain morphology is more uniform, the internal substructure distribution is more continuous, and the overall anisotropy of the microstructure is significantly reduced. The alternating "reaustitization-hot deformation-rapid cooling" process during the two rounds of TMT effectively promotes grain refinement and dislocation reconstruction, causing some high dislocation regions to recover and redistribute in the second round of deformation. After tempering, a relatively uniform dislocation-subgrain structure system is formed inside the material, and a certain number of fine carbides are precipitated, effectively alleviating the local stress concentration problem existing in a single TMT. While maintaining a high strength level, this example shows a significant increase in elongation and impact toughness compared to Example 2, indicating that increasing the number of TMT rounds can restore the toughness of the material to a certain extent, achieving a better match between strength and plasticity.

[0047] Example 4 (using a three-cycle TMT) The alloy composition is the same as in Example 1.

[0048] (1) Batching and smelting: Same as in Example 1.

[0049] (2) Homogenization and forging: Same as in Example 1.

[0050] (3) Multi-cycle thermomechanical treatment (TMT3): The forged sample was first heated to 1100 ℃ and held for 30 min. The first round of hot deformation was carried out at 780 ℃ with a deformation amount of 20%. After hot deformation, it was quickly water-cooled to room temperature. Then, the temperature was raised again to 980 ℃ and held for 20 min. The second round of hot deformation was carried out at 770 ℃ with a deformation amount controlled at 25%. After deformation, it was immediately water-cooled. The temperature was raised again to 900 ℃ and held for 20 min for the third round of hot deformation. The deformation temperature was controlled at 760 ℃ with a deformation amount of 50%. After deformation, it was immediately water-cooled to room temperature.

[0051] (4) Tempering treatment: After the TMT cycle is completed, temper at 740 ℃ for 60 min and air cool.

[0052] Example 4 employs a three-cycle thermomechanical treatment, representing the highest level of TMT (thermomechanical treatment) in this group of examples. Its tempered microstructure is as follows: Figure 7 As shown, the microstructure is highly homogeneous, with further refinement of the grain size, an equivalent grain size of approximately 5 μm, and grain boundary morphology tending towards equiaxed, indicating a significant reduction in microstructural anisotropy. Further SEM and TEM characterization results ( Figure 8 The results showed that after three rounds of TMT, a large number of fine M particles were dispersed and precipitated at the grain boundaries and within the grains during the tempering process. 23 C6-type carbides, these precipitates are discontinuously distributed and effectively pinnate grain boundary migration. Simultaneously, nanoscale MX-type precipitates with a size of approximately 30 nm are observed within the matrix, which interact stably with dislocations and substructures, significantly improving the thermal stability of the microstructure. In this embodiment, the high-density dislocations introduced by multiple rounds of TMT are rationally controlled during tempering, avoiding the risk of embrittlement caused by excessively high dislocation density. This allows the material to maintain high strength at high temperatures, while also achieving excellent impact toughness at room temperature and -20 °C.

[0053] The room temperature (25 °C) and high temperature (650 °C) mechanical properties of the final heat-treated samples of Examples 1, 2, 3 and 4 are shown in Table 1.

[0054] Table 1. Tensile mechanical properties of the examples at room temperature (25 °C) and high temperature (650 °C)

[0055] Table 2. Impact performance at room temperature and -20℃

[0056] As shown in Table 1, Figure 2 and Figure 3 As shown in Examples 2-4, compared with the untreated Example 1, after different numbers of thermomechanical combined treatments (TMT), the alloy showed a significant improvement in strength at both room temperature and high temperature, while the elongation after fracture only fluctuated slightly. Further comparison of different TMT cycles reveals that when the number of TMT cycles gradually increased from 1 (TMT1) to 2 (TMT2) and 3 (TMT3), although the yield strength slightly decreased, the elongation and impact energy showed a continuous upward trend, indicating that increasing the number of TMT cycles helps to restore and improve toughness, achieving a better balance between strength and toughness. Microstructure results ( Figure 4 –7) shows that the NT sample (Example 1) has coarse grains (approximately 50 μm); after a single TMT (Example 2), the grains are significantly refined; after three TMTs, the microstructure becomes significantly more homogeneous and refined, with an equivalent grain size of approximately 5 μm. Regarding precipitates ( Figure 8 After three rounds of TMT, fine and diffuse M can be obtained. 23 C6, distributed along grain boundaries / within grains, effectively "pins" grain boundary migration; under the "three-round TMT + tempering" process, it further induces a large number of nanoscale MX phases (approximately 30 nm in size) to disperse and precipitate, which interact stably with dislocations / substructures, significantly enhancing high-temperature strength and thermal stability of the structure.

[0057] The aforementioned evolution of mechanical properties can be attributed to the following synergistic mechanisms: ① The Hall-Petch strengthening brought about by grain refinement significantly improves yield and tensile strength; ② The high-density dislocations and subgrain structures introduced by multiple rounds of TMT are "locked" after tempering, and together with the dispersed precipitates, they provide dislocation bypass (Orowan) strengthening; ③ Fine, discontinuous M 23 The synergistic pinning of grain boundaries and dislocations by C6 and nano-sized MX effectively suppresses grain growth and high-temperature softening, while avoiding the potential adverse effects of continuous network carbides on toughness; ④ With the increase of TMT cycles, local recovery / partial recrystallization leads to a moderate decrease in dislocation density, explaining the slight decrease in yield strength, but this also reduces the tendency to embrittlement and promotes the improvement of elongation and impact energy. In summary, by rationally designing the TMT cycles and coordinating with tempering to control precipitation, "grain refinement + dispersed MX" was achieved. 23 The multi-dimensional synergistic strengthening / toughening pathway of "C6+nanoMX" enables the alloy of this invention to simultaneously achieve high strength and high toughness under both room temperature and high temperature conditions.

[0058] The embodiments described above merely illustrate specific implementation methods of this application, and while the descriptions are detailed and specific, they should not be construed as limiting the scope of protection of this application. It should be noted that those skilled in the art can make various modifications and improvements without departing from the concept of the technical solution of this application, and these modifications and improvements all fall within the scope of protection of this application.

Claims

1. A high-temperature resistant, low-activation steel, characterized in that, Its chemical composition consists of the following raw materials in the following mass percentages: Cr: 8.0–10.0%, W: 1.0–2.0%, V: 0.2–0.5%, Ta: 0.05–0.2%, Y: 0.005–0.02%, Ce: 0.001–0.005%, B: 0.001–0.005%, C: 0.08–0.16%, N: 0.015–0.035%, Mn: 0.3–0.8%, with the balance being Fe; The steel matrix is ​​tempered martensite, with an original austenite grain equivalent size of less than 10 μm, and contains dispersed MX-type carbides (nitrides) and M... 23 C6 type carbides.

2. The high-temperature resistant, low-activation steel according to claim 1, characterized in that, The mass fractions of highly active elements and impurity elements in the steel are as follows: Ni≤0.02%, Co≤0.01%, Nb≤0.01%, Mo≤0.01%, Cu≤0.01%, Al≤0.05%, Si≤0.10%, P≤0.005%, S≤0.005%, O≤0.005%.

3. A high-temperature resistant, low-activation steel according to claim 1 or 2, characterized in that, The MX-type carbides (nitrides) are mainly Ta and V-based carbides (nitrides) with an average size of 10–60 nm; the M 23 C6 type carbides have an average size of 50–200 nm, are diffusely distributed along the grain interior and grain boundaries, and are not in a continuous network.

4. The high-temperature resistant, low-activation steel according to claim 1, characterized in that, Tensile strength at 650℃ R m ≥400 MPa and yield strength R p0.2 ≥350 MPa, elongation after fracture A ≥20%; at room temperature and / or Impact absorption energy at 20℃ A KV2 ≥150 J.

5. The application of the high-temperature resistant, low-activation steel as described in any one of claims 1 to 4 as a structural material for nuclear fusion reactor cores.

6. The application according to claim 5, characterized in that, Used in the first wall / tritium-producing blanket, divertor and other structures of fusion reactors, as well as their pressure-bearing / heat-conducting components.

7. The method for preparing a high-temperature resistant, low-activation steel according to any one of claims 1 to 4, characterized in that, Includes the following steps: (1) Vacuum induction melting: Melting under a vacuum of ≤1 Pa and casting into steel ingots; (2) Homogenization and forging: After the steel ingot is homogenized by holding at 1100~1250 ℃ for ≥5 h, it is forged by alternating upsetting and drawing at 1100~1200℃ for ≥3 times, with a forging ratio of ≥4 and a final forging temperature of ≥850 ℃; (3) Multi-cycle thermomechanical treatment: First round: Heat the forged billet after step (2) to 1000~1150 ℃ and hold for ≥30 min, then perform hot deformation at 750~800℃ with a deformation amount ≤30%, and then cool it to room temperature by water or spray. Circulating wheel: The billet is reheated to 900-1000 ℃ and held for ≥20 min. It is then hot deformed at 750-800 ℃ with a deformation of 35-60%. After that, it is cooled to room temperature by water or spray. (4) Tempering: Hold at 700-760 ℃ for 30-90 min and then air-cool or air-cool; finally obtain tempered martensite matrix, fine original austenite grains and form dispersed MX and M 23 High-temperature resistant, low-activation steel with C6 precipitates.

8. The method for preparing a high-temperature resistant, low-activation steel according to claim 7, characterized in that, In step (3), each round of hot deformation in the multi-cycle thermomechanical treatment process is carried out by hot compression or hot rolling, and after each round of hot deformation is completed, it is not kept at an isothermal temperature but is directly and rapidly cooled to room temperature.

9. The method for preparing a high-temperature resistant, low-activation steel according to claim 8, characterized in that, The holding time for tempering in step (4) is 45 to 90 minutes.

10. The method for preparing a high-temperature resistant, low-activation steel according to claim 7, characterized in that, In step (3), the cycle wheel is executed two to three times.

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