4.2 K steel with yield strength of 1,500 MPa for nuclear fusion superconducting armor structure and manufacturing method of 4.2 K steel

Abstract

The invention belongs to the technical field of stainless steel, and discloses 4.2 K steel for a nuclear fusion superconducting armor structure with the yield strength of 1500 MPa and a manufacturing method of the 4.2 K steel. The steel comprises the following chemical components: 0.01-0.04% of C, 0.25-0.45% of Si, 5.00-8.00% of Mn, 21-22.5% of Cr, 10-13% of Ni, 2-3% of Mo, 0.35-0.40% of N, 0.10-0.30% of V, 0.05-0.15% of Nb, 0.15-0.25% of Cu, less than or equal to 0.002% of S and less than or equal to 0.02% of P. The top temperature before electroslag billet cogging and upsetting is 1180-1220 DEG C, and the top temperature during drawing is 900-940 DEG C; the top temperature after cogging is 1180-1220 DEG C, the total deformation rate of rough rolling is less than or equal to 60%, the initial rolling temperature of finish rolling is 910-930 DEG C, the total deformation rate of finish rolling is less than or equal to 50%, cooling is performed after rolling, and the re-redness is less than or equal to 400 DEG And the solid solution temperature is 1140-1160 DEG C, and the quenching cooling speed after solid solution is 39.5-40.5 DEG C / s. The yield strength of the 4.2 K steel plate is larger than or equal to 1500 MPa, the ductility is larger than or equal to 30%, the fracture toughness is larger than or equal to 180 MPa.m < 1 / 2 >, and the problems of air holes and edge cracks of high-nitrogen austenitic stainless steel are solved.

Description

Technical Field

[0001] The present invention belongs to the technical field of stainless steel, and specifically relates to a stainless steel plate for a superconducting magnet support structure of a controlled nuclear fusion reactor, and more particularly to a steel for a nuclear fusion superconducting armor structure with a 4.2K yield strength of 1500MPa and a manufacturing method thereof. Background Art

[0002] Compared with the International Thermonuclear Experimental Reactor (ITER), China's new generation fusion engineering test reactor has a superconducting magnet with a maximum magnetic field of 20T and an electromagnetic force of 1400kN / m, which are 54% and 75% higher than the ITER experimental reactor respectively. ITER's general low-temperature structural materials can no longer meet engineering requirements. Therefore, it is urgent to develop a new generation of high-strength and tough stainless steel for nuclear fusion engineering.

[0003] According to the design requirements of key structural components (coil boxes, armor) of China's new generation fusion project, the components are required to meet the coordinated matching of strength and toughness at the same time under the ultra-low temperature and strong magnetic field environment of 4.2K, and strictly avoid the production of α' martensite and ε martensite. Therefore, the common phase transformation strengthening, strain aging strengthening, precipitation strengthening and work hardening of ultra-high strength austenitic stainless steel are not suitable for high strength and toughness stainless steel used in nuclear fusion engineering. Only solid solution strengthening and fine grain strengthening dominated by twin deformation at low temperature can be adopted.

[0004] The nitrogen element in austenitic stainless steel can play a role in solid solution strengthening, significantly improving the yield strength and tensile strength. Combined with the high manganese composition design, it can regulate the stacking fault energy and promote deformation twinning to coordinate plastic deformation, thereby offsetting the problem of reduced low-temperature toughness caused by nitrogen, and greatly improving ultra-low temperature plasticity and toughness. At the same time, manganese and nitrogen elements can improve the stability of austenite. Even if the structural parts are in service at ultra-low temperatures and huge electromagnetic stresses, no α' martensite and ε martensite will be produced, and they will always remain non-magnetic. Therefore, high manganese and nitrogen austenitic stainless steel is an ideal material to ensure the stable operation of fusion reactors.

[0005] However, the production of Fe-Cr-Mn-Mo-N high-manganese and high-nitrogen austenitic stainless steel mainly faces the following problems: (1) When smelted under normal pressure, after adding high concentration nitrogen, local segregation will cause the solidification mode to change from A-type solidification to FA-type solidification. The latter will produce delta ferrite and form "ferrite traps". The solubility of nitrogen will be greatly reduced, causing a large number of subcutaneous bubbles in the ingot, seriously affecting the yield rate. At the same time, the strength of the steel plate is difficult to ensure. (2) Austenitic stainless steel itself has a large linear expansion coefficient and low thermal conductivity. It has a large residual stress during hot rolling. The edge of the rolled piece will increase the tensile stress due to the rapid temperature drop in the late deformation stage. In conjunction with the precipitation phases such as Z phase and Cr2N, it will cause rolling edge cracking. (3) Manganese is an element that easily segregates. The high manganese composition design will lead to a sharp increase in the amount of manganese segregation on the grain boundary. After conventional solid solution treatment, manganese will produce unbalanced segregation at the austenite grain boundary, causing grain boundary brittle fracture and deteriorating low-temperature fracture toughness. The above problems have seriously restricted the industrial promotion of high manganese and nitrogen austenitic stainless steel plates in the field of fusion. Summary of the Invention

[0006] The problems to be solved by the present invention are the porosity caused by local nitrogen segregation in the smelting of high-nitrogen austenitic stainless steel under normal pressure, the edge cracking of steel plates caused by precipitated phases, and the significant decrease in ultra-low temperature plasticity and toughness caused by manganese segregation. In response to the common technical problems in the production of the above-mentioned high-manganese-nitrogen austenitic stainless steel, the present invention provides a 4.2K yield strength 1500MPa-grade steel for nuclear fusion superconducting armor structure and a manufacturing method thereof. Through the composition design, deformation process and heat treatment method of high-manganese-nitrogen austenitic stainless steel, a yield strength of more than 1500MPa at 4.2K temperature and a fracture toughness of more than 130MPa·m are developed. 1 / 2 The high-nitrogen austenitic stainless steel hot-rolled medium and thick plates ensure that the performance indicators of the plates meet the design requirements of my country's latest generation of fusion pilot experimental reactors. This technological achievement can also be applied to superconductors, hydrogen storage and other fields with special requirements for ultra-low temperature toughness and ultra-low temperature strength.

[0007] To achieve the above object, the technical solution adopted by the present invention is as follows:

[0008] One of the technical solutions of the present invention is to provide a 4.2K yield strength 1500MPa-grade steel for nuclear fusion superconducting armor structures, the chemical composition of which is as follows by weight: C: 0.01%-0.04%, Si: 0.25%-0.45%, Mn: 5.00%-8.00%, Cr: 21%-22.5%, Ni: 10%-13%, Mo: 2%-3%, N: 0.35%-0.40%, V: 0.10%-0.30%, Nb: 0.05%-0.15%, Cu: 0.15%-0.25%, S≤0.002%, P≤0.02%, and the remainder being Fe and unavoidable impurities. Among the unavoidable impurities, B≤0.002% and Al≤0.01%. The ratio of Nb to V is 1:1.8-2.2.

[0009] The reasons for adopting the above components are as follows:

[0010] The design mass fraction of Cr is 21%~22.5%. Cr plays a role in solid solution strengthening and can greatly increase the solid solubility of nitrogen. However, in this composition system, if the Cr content exceeds 22.5%, it will promote the formation of σ phase and M 23 The formation of C6 and Cr2N seriously deteriorates low-temperature toughness and hot workability;

[0011] The designed mass fraction of Ni is 10% to 13%, which inhibits the formation of delta ferrite, plays a role in solid solution strengthening, and improves low-temperature toughness;

[0012] The designed mass fraction of Mo is 2% to 3%, which plays a role in solid solution strengthening and forming Laves phase precipitation strengthening without affecting the stability of austenite, thereby improving the ultra-low temperature strength. Further increasing Mo will promote the formation of delta ferrite.

[0013] The design mass fraction of Mn is 5.00%~8.00%. Its function is the same as Ni, which can increase the stability of austenite at low temperatures. In particular, Mn can increase the Neel temperature and prevent the steel from becoming ferromagnetic at low temperatures. In addition, adding 5.00%~8.00% Mn can control the stacking fault energy to be lower than 20mJ / m 2 , which can promote deformation twinning to coordinate plastic deformation and offset the problem of low-temperature toughness reduction caused by nitrogen;

[0014] The designed mass fraction of N is 0.35% to 0.40%. N plays a role in solid solution strengthening and greatly improves the ultra-low temperature strength. Considering that the maximum solid solubility of N under the composition of the alloy of the present invention is about 0.40%, further increasing the N content will cause it to escape in molecular form, forming subcutaneous pores;

[0015] The design mass fraction of V is 0.10%~0.30%. Considering that Nb-containing high nitrogen stainless steel has complex Z phase (Cr3Nb3N and CrNbN) precipitation during hot working, the Z phase is dispersed and precipitated at the grain boundary, deteriorating the hot working performance and causing edge cracks in the steel plate. The composite addition of V and Nb can promote the dispersion and precipitation of Nb in the form of MX and M2X phases before hot working. The small size does not affect the hot working performance, and a large amount of Nb is consumed, thereby inhibiting Z-direction precipitation and improving the hot working performance.

[0016] In addition, the alloy under the composition system of the present invention is controlled at 0.05% to 0.15% Nb and 0.10% to 0.30% V, which can form a chemical formula of Cr 1.75 Nb 0.15 V 0.25 Compared with conventional NbN or VN, the MX phase of N3 significantly increases the ability of microalloying elements to bind nitrogen, significantly improving the "nitrogen fixation" efficiency, allowing nitrogen atoms to dissolve in the steel liquid in the form of compounds during solidification, and dissolve in the alloy matrix in the form of interstitial atoms after solidification, thus avoiding the formation of pores in the cast structure;

[0017] The designed mass fraction of Cu is 0.15%~0.25%, which is used to form ε-Cu, improve the room temperature and low temperature strength, and improve the processing performance of the steel plate;

[0018] B is designed to be ≤ 0.002%. Considering that the segregation of boron impurities in the metal matrix may reduce the material's resistance to radiation swelling and high-temperature strength, affecting the life of key components such as the first wall, controlling the content below 0.002% can effectively prevent its segregation;

[0019] The Al content is designed to be ≤ 0.01% because high-energy neutrons bombarding aluminum atoms will destroy their lattice structure, generate vacancies and interstitial atoms, cause material swelling and embrittlement, and reduce the service life of the components.

[0020] Composition design features: low stacking fault energy is achieved through high manganese composition design, deformation twinning is promoted to coordinate plastic deformation, and the strength is greatly improved by the solid solution strengthening effect of nitrogen to meet the high strength and high plasticity and toughness requirements under ultra-low temperature conditions of 4.2K. By adding niobium and vanadium in combination and controlling their proportions, the second phase composite precipitation behavior of microalloying elements niobium and vanadium in high nitrogen austenitic stainless steel is different from the second phase precipitation behavior of adding niobium or vanadium alone, that is, the chemical formula Cr is formed. 1.75 Nb 0.15 V 0.25Compared with conventional NbN or VN, the MX phase of N3 significantly increases the ability of microalloying elements to bind nitrogen, and the "nitrogen fixation" efficiency is significantly improved. Its precipitation behavior is characterized by starting to precipitate from the liquid phase at 1251°C and dissolving in the austenite at 1131°C, so that nitrogen atoms can be dissolved in the steel liquid in the form of compounds during solidification, and after solidification, they are dissolved in the alloy matrix in the form of interstitial atoms. The nitrogen fixation effect is significantly improved, thereby inhibiting the grain boundary precipitation of Cr2N during rolling, and effectively solving the problems of porosity defects and reduced nitrogen solubility caused by the "ferrite trap" problem of high-nitrogen austenitic stainless steel smelted at normal pressure.

[0021] The second technical solution of the present invention is to provide a method for manufacturing steel for nuclear fusion superconducting armor structure with a yield strength of 4.2K and a strength of 1500MPa, wherein the electroslag billet is opened, wherein the upsetting forging ratio is greater than 4, the drawing forging ratio is greater than 3, and the total forging ratio is greater than 8; the top temperature of the electroslag billet before upsetting deformation is 1180-1220°C, and the insulation is 0.20-0.30min / mm; the top temperature of the drawing section is controlled at 900-940°C, and the insulation is 0.20-0.30min / mm; after forging, the top temperature of the soaking section is controlled at 1180-1220°C, and the insulation is 0.1-0.3 min / mm; full longitudinal rolling is adopted for rolling, the average single-pass reduction rate of rough rolling is ≤25%, the total deformation rate is ≤60%, the starting rolling temperature of finishing rolling is 910~930℃, the average single-pass reduction rate of finishing rolling is ≤15%, the total deformation rate is ≤50%, laminar cooling is adopted after rolling, and the red-return temperature is ≤400℃; the solution treatment temperature is 1140~1160℃, the holding time is 1.1~1.3min / mm, and the quenching cooling rate after solution treatment is controlled at 39.5~40.5℃ / s.

[0022] The above processes and their parameters were chosen because:

[0023] 1. The total forging ratio of the electroslag ingot is greater than 8. The purpose is to increase the deformation of the electroslag ingot core, homogenize the structure, and widen the intermediate billet through upsetting plastic deformation to complete more than 80% of the thickness deformation and reach the width of the finished steel plate. The upsetting forging ratio exceeds 4, which can promote full recrystallization of the edge of the intermediate billet during the widening process, increase the grain boundary area, and effectively suppress the cracks caused by residual stress, tensile stress, etc. to prevent their expansion.

[0024] 2. Before upsetting, the top temperature of the electroslag billet is controlled at 1180-1220°C, with a holding time of 0.20-0.30 min / mm. This ensures uniform heating of the electroslag ingot, achieves high-temperature deformation during upsetting, improves the diffusion capacity of nitrogen in the austenite lattice, reduces local nitrogen segregation, and utilizes the high-temperature softening effect and the solid solution effect of nitrogen to reduce deformation tensile strength, ensuring smooth completion of large upsetting deformation. The second fire drawing process controls the top temperature at 900-940°C with a holding time of 0.20-0.30 min / mm. The purpose is to orient the austenite grains of the intermediate billet along the drawing direction through drawing at 900-940°C, obtaining a fibrous structure and inhibiting the propagation of thermal cracks.

[0025] 3. After forging, the cooling rates of the intermediate billet surface and core differ significantly, resulting in asynchronous thermal expansion and contraction, which generates thermal stress. Furthermore, its poor thermal conductivity further exacerbates the temperature gradient effect. Staying in the preheating section and slowly controlling the heating rate are beneficial for releasing thermal stress and promoting the dispersion and precipitation of the MX and M2X phases. These phases pin the austenite grain boundaries, deformation bands, dislocations, and other areas, thereby pinning the grain boundaries and inhibiting grain coarsening in the intermediate billet. Controlling the top temperature of the soaking section to 1180–1220°C and holding for 0.1–0.3 min / mm2 primarily eliminates residual delta ferrite.

[0026] 5. Full longitudinal rolling is adopted, on the one hand, to reduce the edge extension stress, and on the other hand, it can eliminate the transverse rolling and steel conversion operation. The remaining approximately 20% of the deformation is completed by longitudinal rolling, which greatly shortens the rolling time and reduces the temperature drop of the steel plate edge, ensuring that the edge temperature is above 850°C, reducing the material deformation resistance, reducing the risk of crack initiation caused by local stress concentration at the edge, and solving the problem of edge cracking during rolling.

[0027] 6. When the total deformation rate of rough rolling is greater than 60%, severe plastic deformation may induce the transformation of austenite to martensite. At the same time, the phase transformation volume effect leads to local stress concentration, and microcrack sources are easily formed at grain boundaries or inclusions. When the single-pass reduction rate of rough rolling is greater than 25%, it is easy to cause deformation heat accumulation, local temperature rise, thereby accelerating the dynamic precipitation of Cr2N, etc., leading to cracking. When the total deformation rate of finishing rolling is greater than 50%, it is easy to promote the precipitation of σ phase along the grain boundaries, destroy the grain boundary bonding force, and easily expand into cracks during the finishing rolling process. When the average single-pass reduction rate of finishing rolling is greater than 15%, the accumulated residual stress will exceed the tensile strength of the material and cause cracking.

[0028] 7. The finishing rolling (second rolling) temperature is 910~930℃. This parameter is selected mainly because the 900~950℃ range is the optimal precipitation temperature range for M6C type carbonitrides. The second rolling temperature of 910~930℃ can ensure that M6C can be precipitated in the recrystallized grains during the warming process before the second rolling, thereby further refining the grains during rolling in the non-recrystallized area.

[0029] 8. The steel plate is laminar cooled to a red-return temperature ≤ 400℃ after rolling. The reason is that after rolling deformation, the degree of deformation-accelerated coarsening is related to the second phase precipitation temperature and the order of recovery and recrystallization. During the recovery process, small-angle grain boundaries and substructures promote diffusion, which increases the coarsening kinetic rate. By quickly quenching the steel plate after rolling and making the red-return temperature of the steel plate ≤ 400℃, the recovery from the surface of the steel plate to the core can be suppressed and the atomic diffusion rate of V and Nb can be reduced, thereby reducing the second phase coarsening.

[0030] 9. After hot rolling, the steel plate is subjected to solid solution at 1140-1160°C. Considering that Nb / V is added at a ratio of 1.8-2.2 under the high nitrogen austenite composition system of the present invention, Cr 1.75 Nb 0.15 V 0.25 The MX phase of N3 begins to dissolve in austenite at 1131°C. Setting the solution temperature at 1150°C allows the MX phase to be fully dissolved. The remaining second phases M2X, M6C, etc. of the present invention can also be fully dissolved at this temperature, ensuring that no second phases are precipitated in the solid solution state, thereby significantly improving the ultra-low temperature fracture toughness.

[0031] 10. Considering that conventional high manganese and high nitrogen austenitic stainless steel will change from toughness to brittleness as the temperature decreases, that is, from dimple fracture to grain boundary fracture, the fracture toughness at 4.2K is less than 100MPa·m 1 / 2 The reason is that during conventional solution treatment and heat preservation, a "complex" of vacancies and solute atoms is formed. According to conventional solution treatment, when the cooling rate is around 340℃ / s, the "complex" moves from the intragranular to the grain boundary. As the mass fraction of Mn increases, the segregation of Mn on the grain boundary increases sharply, and eventually unbalanced segregation of Mn occurs at the austenite grain boundary, causing brittle fracture of the grain boundary and seriously deteriorating the low-temperature fracture toughness. By controlling the quenching cooling rate after solution treatment to 39.5~40.5℃ / s, a homogenization process (anti-segregation) of Mn atoms diffusing from the grain boundary to the intragranular can be produced, reducing the segregation of Mn at the grain boundary and improving the low-temperature toughness. Further reducing the cooling rate will lead to carbide precipitation, which will seriously deteriorate the low-temperature toughness. Taking comprehensive consideration, the optimal cooling rate after solution treatment of the present invention is set to 39.5~40.5℃ / s.

[0032] Deformation process design features: The process flow of forging blanking + hot rolling finished steel plates is designed. The intermediate blank is widened through upsetting and large plastic deformation, completing a deformation of more than 80% of the thickness and reaching the width of the finished steel plate. The edge of the intermediate blank is fully recrystallized, the grain boundary area is increased, and cracks caused by residual stress, tensile stress, etc. are effectively suppressed to prevent their expansion; the hot rolling process only performs longitudinal rolling and widening on the rolled piece, which significantly reduces the edge extension stress. Since the longitudinal rolling completes the remaining approximately 20% of the deformation, the rolling time is greatly shortened, the temperature drop at the edge of the steel plate is reduced, and the edge temperature is ensured to be above 850°C, exceeding the precipitation temperature of brittle phases such as Cr2N, avoiding the problem of rolling edge cracking.

[0033] Features of heat treatment process design: By reducing the quenching cooling rate after solution treatment, the Mn atoms can undergo a homogenization process (anti-segregation) of diffusion from the grain boundary to the grain interior, reducing the segregation of Mn at the grain boundary and improving the low-temperature fracture toughness.

[0034] Furthermore, the blanking process adopts the "one upsetting and one drawing" process, which is completed in 2 fires. The total forging ratio is calculated as follows: 总 =y 镦粗 +y 拔长 , where y 镦粗 =H0 / H1;y 拔长 =L1 / L0, H0 is the thickness of the slab before upsetting, mm; H1 is the thickness of the intermediate billet after upsetting, mm; L0 is the length of the slab before drawing, mm; L1 is the length of the intermediate billet after drawing.

[0035] Furthermore, the electroslag billet is upset in the first fire, and the heating system before upsetting is: the electroslag billet is placed in the furnace when the furnace temperature reaches 730~770℃, and the heating rate is controlled at 0.4~0.6℃ / min.

[0036] Furthermore, the electroslag billet is stretched in the second heat. The heating regime before stretching is as follows: the furnace temperature reaches 730-770°C, the intermediate billet is placed, and the heating rate is controlled at 0.4-0.6°C / min. The steel plates produced by this method can be used in the field of ultra-low temperature structural engineering.

[0037] Compared with the prior art, the present invention has the following beneficial effects:

[0038] 1. The present invention limits the range of Nb and V components and limits the addition ratio of the two to 1:1.8~2.2, thereby obtaining a chemical formula of Cr 1.75 Nb 0.15 V 0.25 The MX phase of N3 has a coherent or semi-coherent interface with the high-temperature austenite matrix, which significantly reduces the interfacial energy and promotes its precipitation in the high-temperature austenite matrix, thereby improving the nitrogen fixation ability of the alloy. Even in the case of local nitrogen supersaturation during solidification, no pores will be generated.

[0039] 2. The present invention limits the range of Nb and V components and the addition ratio of the two, so that the formed MX phase can be dispersed and precipitated at high temperature during the heating process of hot working, with a size of less than 30nm, thus not affecting the hot working performance. It can consume a large amount of Nb, thereby inhibiting Z-direction precipitation and improving the hot working performance.

[0040] 3. The present invention improves the grain boundary segregation of Mn by reducing the cooling rate of the steel plate, thereby significantly improving the ultra-low temperature toughness.

[0041] 4. The present invention fully utilizes the solid solution strengthening effect of the alloy through Nb and V microalloying and high N composition design. Combined with the improvement of subsequent smelting, hot working and heat treatment processes, the Nb and V compounds are dispersed and precipitated in the form of MX phase, which inhibits the precipitation of large-sized MX phase, M2X phase, M6C and Z phase, and eliminates the segregation of Mn at the grain boundary. Under the ultra-low temperature condition of 4.2K, its yield strength is ≥1500MPa, tensile strength is ≥1800MPa, elongation is ≥30%, and fracture toughness is ≥180MPa·m 1 / 2 . BRIEF DESCRIPTION OF THE DRAWINGS

[0042] Figure 1 Calibration of solidified tissue transmission morphology and MX phase diffraction spots;

[0043] Figure 2 TEM-EDS image of MX phase. DETAILED DESCRIPTION

[0044] The following examples are only some of the best embodiments of the present invention and do not limit the scope of the aforementioned invention and technical means.

[0045] A 4.2K yield strength, 1500 MPa-grade steel for nuclear fusion superconducting armor structures has the following chemical composition by weight: C: 0.01%-0.04%, Si: 0.25%-0.45%, Mn: 5.00%-8.00%, Cr: 21%-22.5%, Ni: 10%-13%, Mo: 2%-3%, N: 0.35%-0.40%, V: 0.10%-0.30%, Nb: 0.05%-0.15%, Cu: 0.15%-0.25%, S ≤ 0.002%, P ≤ 0.02%, and the remainder being Fe and unavoidable impurities. Among the unavoidable impurities, B ≤ 0.002% and Al ≤ 0.01%. The Nb / V ratio is 1:1.8-2.2. Table 1 shows the composition of each embodiment.

[0046]

[0047] A method for manufacturing steel for nuclear fusion superconducting armor structure with a 4.2K yield strength of 1500 MPa, comprising the following steps (1) to (4):

[0048] (1) According to the stainless steel smelting process, a steel ingot or continuous casting billet with the following mass percentage is obtained: the chemical composition of the stainless steel is as follows, calculated by mass percentage: C: 0.01%~0.04%, Si: 0.25%~0.45%, Mn: 5.00%~8.00%, Cr: 21%~22.5%, Ni: 10%~13%, Mo: 2%~3%, N: 0.35%~0.40%, V: 0.10%~0.30%, Nb: 0.05%~0.15%, Cu: 0.15%~0.25, S≤0.002%; P≤0.02%, B≤0.002%, Al≤0.01%, and the rest are Fe and unavoidable impurities. The steel ingot or continuous casting billet smelting process sequentially includes EAF melting, VOD refining and ESR electroslag remelting.

[0049] (2) The electroslag ingot is forged and blanked.

[0050] (3) The intermediate billet is hot rolled to the target thickness of the steel plate by a thick plate rolling mill.

[0051] (4) The hot-rolled steel plate is subjected to solution treatment.

[0052] In step (2), the electroslag blank is opened by the "one upsetting and one drawing" process, which is completed in 2 fires. The total forging ratio is greater than 8, of which the upsetting forging ratio is greater than 4; the drawing forging ratio is greater than 3. The total forging ratio is calculated as follows: 总 =y 镦粗 +y 拔长 , where y 镦粗 =H0 / H1;y 拔长 = L1 / L0, H0 is the slab thickness before upsetting, mm; H1 is the intermediate billet thickness after upsetting, mm; L0 is the slab length before drawing, mm; L1 is the intermediate billet length after drawing. Table 2 shows the intermediate billet dimensions and forging ratios for the steels of various examples.

[0053]

[0054] In step (2), the electroslag billet is subjected to upsetting in the first fire. The heating system before upsetting is as follows: the electroslag billet is placed in the furnace when the temperature reaches 730-770°C, the heating rate is controlled at 0.4-0.6°C / min, the top temperature is 1180-1220°C, and the insulation is 0.20-0.30 min / mm. In step (2), the electroslag billet is stretched in the second fire. The heating system before stretching is as follows: the intermediate billet is placed in the furnace when the temperature reaches 730-770°C, the heating rate is controlled at 0.4-0.6°C / min, the top temperature is 900-940°C, and the insulation is 0.20-0.30 min / mm. Table 3 shows the forging temperature parameters of the steels of each embodiment.

[0055]

[0056] In step (3), the intermediate billet is maintained at a preheating temperature below 400°C for 1.5 to 2.5 hours. The heating rate in the heating section is controlled at 0.4 to 0.6°C / min. The top temperature in the soaking section is 1180 to 1220°C, and the holding time is 0.15 to 0.25 min / mm. Table 4 shows the rolling heating parameters for each example steel.

[0057]

[0058] In step (3), the rolling strategy adopts full longitudinal rolling, the average single-pass reduction of rough rolling is ≤25%, the total deformation rate is ≤60%, the second rolling temperature is 910-930°C, the average single-pass reduction of finishing rolling is ≤15%, the total deformation rate is ≤50%, the steel plate thickness ranges from 12 to 50 mm, and laminar cooling is adopted after rolling, with the red-hot temperature ≤400°C. Table 5 shows the rolling parameters of each example steel.

[0059]

[0060] In step (4), the solution treatment temperature is 1140-1160°C, the holding time is 1.1-1.3 min / mm, and the quenching cooling rate after solution treatment is controlled at 39.5-40.5°C / s. Table 6 Solution treatment parameters of various example steels.

[0061]

[0062] At 4.2K, the yield strength of the steel plate is ≥1500MPa; the tensile strength is ≥1800MPa; the elongation is ≥30%; and the fracture toughness is ≥130MPa·m 1 / 2 Table 7 shows the mechanical properties of the steels of various examples at 4.2K.

[0063]

[0064] Figure 1 The continuous casting billet produced in Example 1 was cut online at 1200°C into test pieces with a size of 2200*700*50mm. The pieces were immediately water quenched and then machined. The transmission morphology of the resulting solidified structure showed that a large number of precipitated phases with a size of 20~30nm were obtained during the solidification of the component in Example 1. After calibration of the diffraction spots, it was determined that the structure was a body-centered cubic structure with lattice constants of a=2.884 and c=4.131.

[0065] Figure 2 for Figure 1 The TEM-EDS composition analysis results of the precipitated phase show that the precipitated phase contains Cr, V, Nb, and N. The precipitated phase is determined to be MX phase, and its chemical formula is Cr according to the atomic ratio. 1.75 Nb 0.15 V 0.25N3. It is proved that although the lattice difference between the body-centered cubic structure and the face-centered cubic matrix is ​​large, the niobium and vanadium nitride composite precipitation phase of the present invention has a coherent or semi-coherent interface, which significantly reduces the interfacial energy and promotes its precipitation in the high-temperature austenite matrix.

[0066] It should be noted that the above embodiments are only for illustrating the technical concept and features of the present invention and are not intended to limit the present invention. Any equivalent replacement or modification that does not depart from the essence of the present invention shall fall within the scope of protection of the present invention.

Claims

1. A 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structure steel, characterized in that: The chemical composition of steel in mass percentage is: C: 0.01%~0.04%, Si: 0.25%~0.45%, Mn: 5.00%~8.00%, Cr: 21%~22.5%, Ni: 10%~13%, Mo: 2%~3%, N: 0.35%~0.40%, V: 0.10%~0.30%, Nb: 0.05%~0.15%, Cu: 0.15%~0.25%, S≤0.002%, P≤0.02%, and the rest are Fe and unavoidable impurities.

2. The 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 1, characterized in that: Among the inevitable impurities, B ≤ 0.002%, Al ≤ 0.01%.

3. The 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 1, characterized in that: Nb / V is 1:1.8~2.

2.

4. The 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 1, characterized in that: The thickness of the steel plate is 12~50mm. Under the ultra-low temperature condition of 4.2K, the yield strength of the steel is ≥1500MPa, the tensile strength is ≥1800MPa, the elongation is ≥30%, and the fracture toughness is ≥180MPa·m 1 / 2 .

5. A method for manufacturing the 4.2K yield strength 1500 MPa grade nuclear fusion superconducting armor structure steel according to any one of claims 1 to 4, characterized in that: The electroslag billet is opened, wherein the upsetting forging ratio is greater than 4, the drawing forging ratio is greater than 3, and the total forging ratio is greater than 8; the top temperature of the electroslag billet before upsetting deformation is 1180~1220℃, and the temperature is kept at 0.20~0.30min / mm; the top temperature of the drawing is controlled at 900~940℃, and the temperature is kept at 0.20~0.30min / mm; after forging, the top temperature of the soaking section is controlled at 1180~1220℃, and the temperature is kept at 0.1~0.3min / mm; the rolling adopts full longitudinal Rolling, the average single-pass reduction rate of rough rolling is ≤25%, the total deformation rate is ≤60%, the finishing rolling start temperature is 910~930℃, the average single-pass reduction rate of finishing rolling is ≤15%, the total deformation rate is ≤50%, laminar cooling is adopted after rolling, and the red-return temperature is ≤400℃; the solution treatment temperature is 1140~1160℃, the holding time is 1.1~1.3min / mm, and the quenching cooling rate after solution treatment is controlled at 39.5~40.5℃ / s.

6. The method for manufacturing a 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 5, characterized in that: The blanking process adopts "one upsetting and one drawing" process, which is completed in 2 fires. The total forging ratio is calculated as follows: 总 =y 镦粗 +y 拔长 , where y 镦粗 =H0 / H1;y 拔长 =L1 / L0, H0 is the thickness of the slab before upsetting, mm; H1 is the thickness of the intermediate billet after upsetting, mm; L0 is the length of the slab before drawing, mm; L1 is the length of the intermediate billet after drawing.

7. The method for manufacturing a 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 6, characterized in that: The electroslag billet is upset in the first fire. The heating system before upsetting is: the electroslag billet is placed when the furnace temperature reaches 730~770℃, and the heating rate is controlled at 0.4~0.6℃ / min.

8. The method for manufacturing a 4.2K yield strength 1500MPa grade nuclear fusion superconducting armor structural steel according to claim 6, characterized in that: The electroslag billet is stretched in the second fire. The heating system before stretching is: the furnace temperature reaches 730~770℃, the intermediate billet is placed, and the heating rate is controlled at 0.4~0.6℃ / min.

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