4.2k yield strength 1500mpa class nuclear fusion superconducting armor structure steel and method of manufacturing the same

Abstract

The application belongs to the technical field of stainless steel, and discloses a 4.2K yield strength 1500MPa-grade nuclear fusion superconducting armor structure steel and a manufacturing method thereof. The chemical composition of the steel is as follows: 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, and P ≤0.02. The electric slag blank is opened, the top temperature before upsetting is 1180-1220 DEG C, and the top temperature during elongation is 900-940 DEG C; the top temperature after opening is 1180-1220 DEG C, the total deformation rate of rough rolling is ≤60%, the opening rolling temperature of finish rolling is 910-930 DEG C, the total deformation rate of finish rolling is ≤50%, the steel is cooled after rolling, the red temperature is ≤400 DEG C, 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 4.2K steel plate is ≥1500MPa, the elongation rate is ≥30%, and the fracture toughness is ≥180MPa·m 1 / 2 , and the problems of porosity and edge cracking of high-nitrogen austenitic stainless steel are solved.

Description

TECHNICAL FIELD

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

[0002] Compared with the International Thermonuclear Experimental Reactor (ITER), the superconducting magnet of the China's new generation of fusion experimental reactor has a maximum magnetic field of 20T and an electromagnetic force of 1400kN / m, which are 54% and 75% higher than those of the ITER experimental reactor, respectively. The ITER general low-temperature structural material cannot meet the engineering requirements, and therefore it is urgent to develop a new generation of high strength and toughness stainless steel for nuclear fusion engineering.

[0003] According to the design requirements of the key structural components (coil box and armor) of China's new generation of fusion engineering, the components are required to meet the matching of strength and toughness under the condition of 4.2K ultra-low temperature and strong magnetic field, and to strictly avoid the generation of alpha prime martensite and epsilon martensite, and therefore the common phase transformation strengthening, strain aging strengthening, precipitation strengthening and work hardening of ultra-high strength austenitic stainless steel cannot be applied to the high strength and toughness stainless steel for nuclear fusion engineering, and only solid solution strengthening and fine grain strengthening dominated by twinning deformation at low temperature can be adopted.

[0004] The nitrogen element in the austenitic stainless steel can play a solid solution strengthening role, significantly improve the yield strength and tensile strength, cooperate with the high manganese component design, control the dislocation energy, promote the deformation twinning to coordinate the plastic deformation, thereby offsetting the problem of low-temperature toughness reduction caused by nitrogen, greatly improving the ultra-low temperature plasticity and toughness, and the manganese and nitrogen elements can both improve the stability of austenite, so that the alpha prime martensite and epsilon martensite will not be generated even if the structural component is served under the action of ultra-low temperature and huge electromagnetic stress, and the non-magnetic property is always maintained. Therefore, the high manganese and nitrogen austenitic stainless steel is an ideal material for ensuring the stable operation of the fusion reactor.

[0005] However, the production of Fe-Cr-Mn-Mo-N series high manganese and high nitrogen austenitic stainless steel mainly faces the following problems: (1) Smelting under normal pressure, after the addition of high concentration of nitrogen, local segregation will lead to a change in the solidification mode, from A-type solidification to FA-type solidification. The latter will produce δ ferrite, forming a "ferrite trap". The solubility of nitrogen will be greatly reduced, causing a large number of subcutaneous bubbles to be generated in the billet, which will seriously affect the yield. At the same time, the strength of the steel plate is difficult to guarantee; (2) Austenitic stainless steel itself has a large coefficient of linear expansion and low thermal conductivity. During hot rolling, there is a large residual stress. In the later stage of deformation of the rolled piece, the edge will be aggravated by the rapid temperature drop, which will lead to rolling edge cracks in conjunction with Z phase, Cr2N and other precipitated phases; (3) Manganese is an element that is easy to segregate. The high manganese composition design will lead to a sharp increase in the amount of manganese segregation at the grain boundaries. After conventional solid solution treatment, manganese will be unbalanced segregated at the austenite grain boundaries, causing grain boundary brittle fracture and deteriorating the low temperature fracture toughness. The above problems seriously limit the industrial promotion of high manganese and nitrogen austenitic stainless steel plates in the field of fusion. Summary of the Invention

[0006] The problems this invention aims to solve are: porosity caused by local nitrogen segregation in the production of high-nitrogen austenitic stainless steel under normal pressure; edge cracking during steel plate rolling caused by precipitated phases; and a significant decrease in ductility and toughness at ultra-low temperatures due to manganese segregation. Addressing these common technical problems in the production of high-manganese-nitrogen austenitic stainless steel, this invention provides a 4.2K yield strength of 1500MPa for nuclear fusion superconducting armor structures and its manufacturing method. Through compositional design, deformation processes, and heat treatment methods for high-manganese-nitrogen austenitic stainless steel, a yield strength exceeding 1500MPa and a fracture toughness exceeding 130MPa·m at 4.2K are achieved. 1 / 2 The high-nitrogen austenitic stainless steel hot-rolled medium-thick plates ensure that the various 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 fields such as superconductivity and hydrogen storage where there are special requirements for ultra-low temperature toughness and ultra-low temperature strength.

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

[0008] One of the technical solutions of this invention is to provide a 4.2K steel with a yield strength of 1500MPa for nuclear fusion superconducting armor structures. The chemical composition by mass percentage is as follows: 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%, with 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.

[0009] The reasons for using the above-mentioned components are as follows:

[0010] The design incorporates Cr at a mass fraction of 21%–22.5%. Cr acts as a solid solution enhancer and significantly increases the solid solubility of nitrogen. However, in this composition system, a Cr content exceeding 22.5% will promote the formation of the σ phase and the M phase. 23 The formation of C6 and Cr2N severely deteriorates low-temperature toughness and hot workability;

[0011] The design incorporates Ni with a mass fraction of 10%~13%, which helps to suppress the formation of δ-ferrite and provides solid solution strengthening, thereby improving low-temperature toughness.

[0012] The design of Mo with a mass fraction of 2%~3% can play the role of solid solution strengthening and Laves phase precipitation strengthening without affecting the stability of austenite, thereby improving the ultra-low temperature strength. Further increasing the Mo content will promote the formation of δ ferrite.

[0013] Adding 5.00%~8.00% Mn by mass has the same effect as Ni in increasing the stability of austenite at low temperatures. In particular, Mn can raise the Nell temperature, preventing the steel from exhibiting ferromagnetism at low temperatures. In addition, adding 5.00%~8.00% Mn can control the stacking fault energy to below 20 mJ / m 2 This can promote deformation twinning to coordinate plastic deformation and offset the problem of nitrogen-induced low-temperature toughness reduction;

[0014] The design incorporates 0.35% to 0.40% N by mass, which acts as a solid solution strengthener and greatly enhances the strength at ultra-low temperatures. Considering that the maximum solid solubility of N in the alloy composition of this invention is around 0.40%, further increasing the N content will cause it to escape in molecular form, forming subcutaneous pores.

[0015] The design uses a V mass fraction of 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 disperses at the grain boundaries, which deteriorates the hot working performance and causes edge cracks in the steel plate. By adding V and Nb in combination, Nb can be promoted to disperse and precipitate in the form of MX and M2X phases before hot working. The size is small and does not affect the hot working performance. It consumes a large amount of Nb, thereby inhibiting Z-phase precipitation and improving the hot working performance.

[0016] Furthermore, the alloy in the composition system of this invention, with Nb controlled at 0.05%~0.15% and V at 0.10%~0.30%, can form an alloy with the chemical formula Cr. 1.75 Nb 0.15 V 0.25 The MX phase of N3 significantly increases the ability of microalloying elements to bind nitrogen compared with conventional NbN or VN, and the "nitrogen fixation" efficiency is significantly improved. This allows nitrogen atoms to be dissolved in the molten steel in the form of compounds during solidification, and to be dissolved in the alloy matrix in the form of interstitial atoms after solidification, thus avoiding the formation of pores in the as-cast structure.

[0017] The design incorporates 0.15% to 0.25% Cu by mass to form ε-Cu, thereby improving room temperature and low temperature strength and enhancing the processing performance of the steel sheet.

[0018] The design B ≤ 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, and affect the lifespan of key components such as the first wall, controlling the content below 0.002% can effectively prevent its segregation.

[0019] The design incorporates Al ≤ 0.01%, taking into account that when high-energy neutrons bombard aluminum atoms, they will disrupt the crystal structure, creating vacancies and interstitial atoms, causing material swelling and embrittlement, and reducing the service life of components.

[0020] Compositional Design Features: A high-manganese composition design achieves low stacking fault energy, promoting deformation twinning to coordinate plastic deformation. The solid solution strengthening effect of nitrogen significantly improves strength, meeting the requirements of high strength and high ductility and toughness under ultra-low temperature conditions of 4.2K. By compositely adding niobium and vanadium and controlling their proportions, the unique precipitation behavior of niobium and vanadium as microalloying elements in the second phase of high-nitrogen austenitic stainless steel is utilized, differing from the precipitation behavior of adding niobium or vanadium alone. This results in the formation of a second phase with the chemical formula Cr... 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 precipitation from the liquid phase at 1251℃ and dissolution in austenite at 1131℃. This allows nitrogen atoms to be dissolved in the molten steel in the form of compounds during solidification, and then dissolved in the alloy matrix as interstitial atoms after solidification. The nitrogen fixation effect is significantly improved, thereby inhibiting the precipitation of Cr2N at grain boundaries during rolling. This effectively solves the problems of porosity defects and reduced nitrogen solubility caused by the "ferrite trap" problem in high-nitrogen austenitic stainless steel smelting under normal pressure.

[0021] The second technical solution of this invention provides a method for manufacturing 4.2K steel with a yield strength of 1500MPa for nuclear fusion superconducting armor structures. The method involves preparing an electroslag billet, wherein the upsetting forging ratio is >4, the drawing forging ratio is >3, and the total forging ratio is >8. Before upsetting, the top temperature of the electroslag billet is 1180~1220℃, held for 0.20~0.30 min / mm. During drawing, the top temperature is controlled at 900~940℃, held for 0.20~0.30 min / mm. After forging, the top temperature of the soaking zone is controlled at 1180~1220℃, held for 0.1~0.3 min / mm. The rolling process uses full longitudinal rolling. The average single-pass reduction rate in roughing is ≤25%, and the total deformation rate is ≤60%. The starting rolling temperature in finishing is 910~930℃, the average single-pass reduction rate in finishing is ≤15%, and the total deformation rate is ≤50%. Laminar flow cooling is used after rolling, and the reddening 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 procedures and parameters were chosen because:

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

[0024] 2. Controlling the top temperature of the electroslag ingot before upsetting deformation to 1180~1220℃ and holding it for 0.20~0.30 min / mm ensures uniform heating of the electroslag ingot, achieves high-temperature deformation during upsetting, improves the diffusion ability 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 the tensile strength during deformation, ensuring the smooth completion of large upsetting deformation. The second drawing temperature is controlled at 900~940℃ and held for 0.20~0.30 min / mm. The purpose of drawing at 900~940℃ is to orient the austenite grains in the intermediate billet along the drawing direction, resulting in a fibrous structure and inhibiting hot crack propagation.

[0025] 3. After forging, the cooling rate difference between the surface and core of the intermediate billet is significant, resulting in asynchronous thermal expansion and contraction and the generation of thermal stress. In addition, its poor thermal conductivity further exacerbates the temperature gradient effect. By holding the billet in the preheating section and slowly controlling the heating rate, it is beneficial to release thermal stress and promote the precipitation of MX and M2X phases in a dispersed manner. These phases pin the billets in austenite grain boundaries, deformation bands, dislocations, etc., thus pinning the grain boundaries and inhibiting the coarsening of the intermediate billet grain size. Controlling the top temperature of the soaking section to 1180~1220℃ and holding it for 0.1~0.3 min / mm mainly serves to eliminate residual δ-ferrite.

[0026] 5. The use of full longitudinal rolling is intended to reduce edge extension stress and eliminate the need for transverse rolling to steel transfer. The remaining approximately 20% of deformation is completed by longitudinal rolling, which greatly shortens the rolling time, reduces the temperature drop at the edge of the steel plate, ensures that the edge temperature is above 850℃, reduces the material's deformation resistance, reduces the risk of crack initiation caused by local stress concentration at the edge, and solves the problem of edge cracking during rolling.

[0027] 6. When the total deformation rate of rough rolling is >60%, the intense 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 initiation is easily formed at grain boundaries or inclusions. When the single-pass reduction rate of rough rolling is >25%, it is easy to cause the accumulation of deformation heat, and the local temperature rises, thereby accelerating the dynamic precipitation of Cr2N and other phases, leading to cracking. When the total deformation rate of finish rolling is >50%, it is easy to promote the precipitation of σ phase along grain boundaries, destroying the grain boundary bonding force. During the finish rolling process, it is easy to extend into cracks. When the average single-pass reduction rate of finish rolling is >15%, the accumulated residual stress will exceed the tensile strength of the material, causing cracking.

[0028] 7. The finishing rolling start (second start) 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 start temperature of 910~930℃ can ensure that M6C can be precipitated in the recrystallized grains during the preheating process before the second start, thereby further refining the grains during the rolling of the non-recrystallized area.

[0029] 8. The steel plate is laminar cooled to a red-hot temperature of ≤400℃ after rolling. The reason is that after rolling deformation, the degree of coarsening due to deformation acceleration is related to the precipitation temperature of the second phase 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 rapidly water quenching the steel plate after rolling and keeping the red-hot temperature of the steel plate surface ≤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 alleviating the coarsening of the second phase.

[0030] 9. After hot rolling, the steel plate is solution treated at 1140~1160℃. Considering that Nb / V is added at a ratio of 1:1.8~2.2 in the high-nitrogen austenite composition system of this invention, Cr is obtained. 1.75 Nb 0.15 V 0.25 The MX phase of N3 begins to dissolve in austenite at 1131℃. Setting a solid solution temperature of 1150℃ allows the MX phase to be fully dissolved. The other second phases of this invention, such as M2X and M6C, can also be fully dissolved at this temperature, ensuring that no second phase precipitates in the solid solution state, thereby significantly improving the ultra-low temperature fracture toughness.

[0031] 10. Considering that conventional high-manganese, high-nitrogen austenitic stainless steel undergoes a transition from ductile to brittle as the temperature decreases, i.e., from dimple fracture to grain boundary fracture, the fracture toughness at 4.2K is less than 100 MPa·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 interior of the grain to the grain boundary. As the mass fraction of Mn increases, the amount of Mn segregation at the grain boundary increases sharply, eventually causing unbalanced segregation of Mn at the austenite grain boundary, resulting in brittle fracture of the grain boundary and severely deteriorating the low-temperature fracture toughness. By controlling the quenching cooling rate after solution treatment to 39.5~40.5℃ / s, a homogenization process (reverse segregation) can be achieved, which allows Mn atoms to diffuse from the grain boundary to the interior of the grain, reducing the amount of Mn segregation at the grain boundary and improving the low-temperature toughness. If the cooling rate is further reduced, it will lead to the precipitation of carbides, which will severely deteriorate the low-temperature toughness. Considering all factors, the optimal cooling rate after solution treatment in this invention is set to 39.5~40.5℃ / s.

[0032] Features of the deformation process design: The design incorporates a forging billet opening process followed by hot rolling of the finished steel plate. Upsetting and large plastic deformation widen the intermediate billet, achieving over 80% of the thickness deformation and reaching the width of the finished steel plate. This allows for sufficient recrystallization at the edges of the intermediate billet, increasing the grain boundary area and effectively suppressing cracks caused by residual stress and tensile stress, preventing their propagation. The hot rolling process only widens the workpiece longitudinally, significantly reducing edge extension stress. Since longitudinal rolling completes approximately 20% of the remaining deformation, it greatly shortens the rolling time, reduces the temperature drop at the steel plate edges, and ensures that the edge temperature remains above 850℃, exceeding the precipitation temperature of brittle phases such as Cr2N, thus avoiding edge cracking during rolling.

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

[0034] Furthermore, the billet is forged using a "one upsetting, one drawing" process, completed in two passes. The total forging ratio is calculated using the formula: y 总 =y 镦粗 +y 拔长 , where y 镦粗 =H0 / H1; y 拔长 =L1 / L0, where H0 is the slab thickness before upsetting (mm); H1 is the intermediate slab thickness after upsetting (mm); L0 is the slab length before drawing (mm); and L1 is the intermediate slab length after drawing.

[0035] Furthermore, the electroslag billet is upset in the first furnace. The heating regime before upset is as follows: 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 undergoes a second drawing process. The heating regime before drawing is as follows: the intermediate billet is placed in the furnace at a temperature of 730~770℃, and the heating rate is controlled at 0.4~0.6℃ / min. The steel plate produced by this method can be applied 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. This invention, by limiting the range of Nb and V components and limiting their addition ratio to 1:1.8~2.2, yields a compound with the chemical formula 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 austenitic matrix, which significantly reduces the interfacial energy and promotes its precipitation in the high-temperature austenitic matrix, thereby improving the nitrogen fixation ability of the alloy. Even under local nitrogen supersaturation during solidification, no porosity will be generated.

[0039] 2. By limiting the range of Nb and V components and the ratio of their addition, the present invention forms an MX phase that can be dispersed and precipitated at high temperature during the heating process of hot working. The phase has a size of less than 30 nm, so it does not affect the hot working performance. It can consume a large amount of Nb, thereby suppressing Z-axis precipitation and improving the hot working performance.

[0040] 3. This invention improves the grain boundary segregation phenomenon of Mn by reducing the cooling rate of the steel plate, and significantly improves the ultra-low temperature toughness.

[0041] 4. This invention, through Nb and V microalloying and high N composition design, fully utilizes the solid solution strengthening effect of the alloy. Combined with improvements in subsequent smelting, hot working, and heat treatment processes, it ensures that Nb and V compounds precipitate in the form of the MX phase, suppressing the precipitation of large-size MX phases, M2X phases, M6C, and Z phases. Simultaneously, it eliminates Mn segregation at grain boundaries. At an ultra-low temperature of 4.2 K, its yield strength ≥1500 MPa, tensile strength ≥1800 MPa, elongation ≥30%, and fracture toughness ≥180 MPa·m. 1 / 2 . Attached Figure Description

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

[0043] Figure 2 This is a TEM-EDS image of the MX phase. Detailed Implementation

[0044] The following embodiments are merely some preferred implementations of the present invention and do not limit the scope and technical means of the invention in any way.

[0045] A 4.2K steel with a yield strength of 1500MPa is used for nuclear fusion superconducting armor structures. Its chemical composition (by mass percentage) is as follows: 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%, with 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 involved in each embodiment.

[0046]

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

[0048] (1) The following mass percentages of steel ingots or continuous casting billets are obtained according to the stainless steel smelting process: The chemical composition of the stainless steel is as follows 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%, with the remainder being Fe and unavoidable impurities. The steel ingot or continuous casting billet smelting process sequentially undergoes EAF melting, VOD refining, and ESR electroslag remelting.

[0049] (2) Electroslag ingots are forged and blanked.

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

[0051] (4) Hot-rolled steel plates are subjected to solution treatment.

[0052] In step (2), the electroslag billet is prepared using a "one upsetting and one drawing" process, completed in two passes, with a total forging ratio > 8, including an upsetting forging ratio > 4 and a drawing forging ratio > 3. The formula for calculating the total forging ratio is: y 总 =y 镦粗 +y 拔长 , where y 镦粗 =H0 / H1; y 拔长 =L1 / L0, where H0 is the slab thickness before upsetting (mm); H1 is the intermediate slab thickness after upsetting (mm); L0 is the slab length before drawing (mm); and L1 is the intermediate slab length after drawing. Table 2 shows the intermediate slab dimensions and forging ratios for the steels in each embodiment.

[0053]

[0054] In step (2), the electroslag billet is upset in the first heat. The heating regime before upset is as follows: the furnace temperature reaches 730~770℃, the electroslag billet is placed in, the heating rate is controlled at 0.4~0.6℃ / min, the top temperature is 1180~1220℃, and the holding time is 0.20~0.30min / mm. In step (2), the electroslag billet is drawn in the second heat. The heating regime before drawing is as follows: the furnace temperature reaches 730~770℃, the intermediate billet is placed in, the heating rate is controlled at 0.4~0.6℃ / min, the top temperature is 900~940℃, and the holding time is 0.20~0.30min / mm. Table 3 shows the forging temperature parameters of the steel in each embodiment.

[0055]

[0056] In step (3), the intermediate billet is kept at a temperature below 400℃ in the preheating section for 1.5~2.5h, the heating rate in the heating section is controlled at 0.4~0.6℃ / min, the top temperature in the soaking section is 1180~1220℃, and the holding time is 0.15~0.25 min / mm. Table 4 shows the rolling heating parameters of the steel in each embodiment.

[0057]

[0058] In step (3), the rolling strategy adopts full longitudinal rolling, with an average single-pass reduction rate of ≤25% and a total deformation rate of ≤60% in roughing, and a second-pass temperature of 910~930℃. In finishing, the average single-pass reduction rate is ≤15% and the total deformation rate is ≤50%. The steel plate thickness ranges from 12 to 50 mm. Laminar flow cooling is used after rolling, and the red-hot temperature is ≤400℃. Table 5 shows the rolling parameters of the steel in each embodiment.

[0059]

[0060] In step (4), 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. Table 6 shows the solution treatment parameters of the steels in each embodiment.

[0061]

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

[0063]

[0064] Figure 1 The continuous casting billet produced in Example 1 was online flame-cut at 1200℃ to a test piece with dimensions of 2200*700*50mm. After immediate water quenching, it was machined. The transmission morphology of the solidified structure obtained at the end showed that the composition of Example 1 obtained a large number of 20~30nm precipitates during solidification. After diffraction spot calibration, it was determined to be a body-centered cubic structure with lattice constants of a=2.884 and c=4.131.

[0065] Figure 2 for Figure 1 TEM-EDS compositional analysis of the precipitated phase showed that it contained Cr, V, Nb, and N, identifying the precipitated phase as the MX phase. Based on atomic ratio, its chemical formula is Cr. 1.75 Nb 0.15 V 0.25N3. It is demonstrated that although the lattice difference between the body-centered cubic structure and the face-centered cubic matrix is ​​large, the niobium-vanadium nitride composite precipitate of the present invention has a coherent or semi-coherent interface, which significantly reduces the interface energy and promotes its precipitation in the high-temperature austenitic matrix.

[0066] It is hereby 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 substitutions or modifications made without departing from the essence of the present invention fall within the protection scope of the present invention.

Claims

1. A type of steel for nuclear fusion superconducting armor structures with a yield strength of 1500 MPa (4.2K), characterized in that, The chemical composition of the steel (by mass percentage) is as follows: 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%, with the remainder being Fe and unavoidable impurities. The steel is manufactured by roughing an electroslag billet, with an upsetting forging ratio >4, a drawing forging ratio >3, and a total forging ratio >8. The top temperature of the electroslag billet before upsetting deformation is 1180~1220℃, and it is held at this temperature. 0.20~0.30 min / mm; drawing temperature controlled at 900~940℃, holding time 0.20~0.30 min / mm; after forging, the top temperature of the soaking zone controlled at 1180~1220℃, holding time 0.1~0.3 min / mm; rolling adopts full longitudinal rolling, roughing average single-pass reduction rate ≤25%, total deformation rate ≤60%, finishing rolling start temperature 910~930℃, finishing average single-pass reduction rate ≤15%, total deformation rate ≤50%, laminar flow cooling is adopted after rolling, and the reddening temperature ≦400℃; solution treatment temperature is 1140~1160℃, holding time is 1.1~1.3 min / mm, and quenching cooling rate after solution treatment is controlled at 39.5~40.5℃ / s.

2. The steel for nuclear fusion superconducting armor structures with a yield strength of 1500 MPa (4.2K) as described in claim 1, characterized in that, Unavoidable impurities include B ≤ 0.002% and Al ≤ 0.01%.

3. The steel for nuclear fusion superconducting armor structure with a yield strength of 1500 MPa (4.2K) as described in claim 1, characterized in that, The Nb / V ratio is 1:1.8~2.

2.

4. The steel for nuclear fusion superconducting armor structure with a yield strength of 1500 MPa (4.2K) as described in claim 1, characterized in that, For steel plates with a thickness of 12~50mm, under ultra-low temperature conditions 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 ≥130MPa·m. 1 / 2 .

5. A method for manufacturing the 4.2K yield strength 1500MPa grade steel for nuclear fusion superconducting armor structures as described in any one of claims 1 to 4, characterized in that, The electroslag billet is prepared by forging, with an upsetting forging ratio >4, a drawing forging ratio >3, and a total forging ratio >8. Before upsetting, the top temperature of the electroslag billet is 1180~1220℃, held for 0.20~0.30 min / mm. During drawing, the top temperature is controlled at 900~940℃, held for 0.20~0.30 min / mm. After forging, the top temperature of the soaking zone is controlled at 1180~1220℃, held for 0.1~0.3 min / mm. Rolling is performed using a full longitudinal rolling process. For roughing, the average single-pass reduction rate is ≤25% and the total deformation rate is ≤60%. For finishing, the initial rolling temperature is 910~930℃, the average single-pass reduction rate is ≤15%, and the total deformation rate is ≤50%. Laminar flow cooling is adopted after rolling, and the reddening 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 steel for nuclear fusion superconducting armor structure according to claim 5, characterized in that, The billet is prepared using a "one upsetting and one drawing" process, completed in two heats. The total forging ratio is calculated as follows: y_total = y_upsetting + y_drawing, where y_upsetting = H0 / H1; y_drawing = L1 / L0, H0 is the thickness of the billet before upsetting (mm); H1 is the thickness of the intermediate billet after upsetting (mm); L0 is the length of the billet before drawing (mm); and L1 is the length of the intermediate billet after drawing.

7. The method for manufacturing a 4.2K yield strength (1500MPa) grade steel for nuclear fusion superconducting armor structure according to claim 6, characterized in that, The electroslag billet is upset in the first furnace. The heating regime before upseting is as follows: the electroslag billet is put into the furnace 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 steel for nuclear fusion superconducting armor structure according to claim 6, characterized in that, The electroslag billet is drawn in the second heating process. The heating regime before drawing is as follows: the furnace temperature reaches 730~770℃ and the intermediate billet is placed in, and the heating rate is controlled at 0.4~0.6℃ / min.

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