Nanometer high-purity high-thermal-stability hot forging die steel and preparation method thereof

Abstract

The application discloses a kind of nano high purity high thermal stability hot forging die steel and preparation method thereof, belong to metal material engineering and die steel technical field.The die steel is according to percentage by mass composition: C 0.38~0.40%, Si 0.20~0.40%, Mn 0.30~0.50%, P≤0.015%, S≤0.002%, Cr 4.80~5.20%, Mo 2.90~3.20%, V 0.40~0.60%, Ni≤0.25%, Cu≤0.10%, every ton steel liquid adopts 6m nano alloy cored wire and is added, the balance is Fe and inevitable impurities.Preparation method adopts medium frequency melting, AOD, LF, VD, die casting, electroslag remelting, forging breakdown, precision forging and spheroidizing annealing ultra-pure short process, so that inclusions and harmful elements in steel are significantly reduced, carbide is fine and dispersed, and grain boundary is effectively pinned.

Description

Technical Field

[0001] This invention belongs to the field of metal materials engineering and mold steel technology, specifically relating to a nano-high purity, high thermal stability hot forging die steel and its preparation method. Background Technology

[0002] Hot forging dies are core tooling in the field of high-end equipment manufacturing. During their service, they are subjected to high temperature and high pressure, repeated hot and cold cycles and impact loads for a long time. Their failure modes are mainly thermal fatigue cracking, high temperature collapse deformation and cracking. Therefore, the die steel is required to have high purity, excellent quenching and tempering stability, thermal fatigue resistance and impact toughness.

[0003] The current mainstream H13 type hot work die steel has three significant shortcomings: First, it has low precision in composition control and high content of impurity elements and gases. Sulfides and oxide inclusions are prone to becoming fatigue crack initiation sites, which restricts the life of the die. Second, it has insufficient resistance to tempering softening at high temperatures. When it is used for a long time at 500~600℃, its hardness drops rapidly, and the die is prone to collapse and deformation. Third, it has limited uniformity of microstructure. Under thermal fatigue cycles, crack initiation and propagation are fast, and the life of the die is difficult to meet the requirements under the conditions of high-tonnage precision hot forging.

[0004] In existing technologies, the matrix is ​​often strengthened by increasing the content of Cr, Mo, and V alloys. However, increasing the alloy content can exacerbate carbide segregation and reduce toughness. Some studies have also attempted to add nanoparticles to refine the microstructure. However, nanoparticles tend to agglomerate in molten steel, have poor dispersibility, and the pinning effect of zero-dimensional particles on grain boundaries is limited, resulting in insufficient improvement in quenching and tempering stability and thermal fatigue performance.

[0005] Patent application CN117925970A discloses a method for preparing hot forging die steel. By controlling the vanadium-carbon ratio (V / C=2.0~2.5), chromium-molybdenum ratio (Cr / Mo=3.0~3.5), and the total amount of Cr, Mo, and V in the steel to be not less than 7.0%, and by adopting a pretreatment process of normalizing, multiple solution treatments, and spheroidizing annealing, the high-temperature wear resistance of the die steel is improved without changing the main composition system. However, this technical solution still has the following limitations: First, the molybdenum content is only 1.50%~1.80%, the precipitation of the secondary hardening phase and high-temperature stability are limited, and the resistance to tempering softening during long-term service above 600℃ is still insufficient, making it difficult to adapt to hot forging conditions at higher temperatures; Second, the in-situ strengthening mechanism of nanophase is not introduced, and it is impossible to refine the microstructure and improve thermal fatigue resistance simultaneously through heterogeneous nucleation and grain boundary pinning; Third, the pretreatment path of two solid solution treatments and two spheroidizing annealing treatments has a long process flow, high production cycle and energy consumption, which limits the efficiency of industrial mass production; Fourth, the content of harmful elements such as S and P and oxygen and nitrogen gases is not strictly controlled, and the risk of thermal fatigue cracks induced by inclusions has not been eliminated at the root, and there is still room for improvement in the transverse toughness and long-life service reliability of large-section dies.

[0006] Therefore, developing a nano-reinforced hot forging die steel with precise and controllable composition, high purity, and excellent quenching and tempering temperature stability and thermal fatigue performance is a technical problem that urgently needs to be solved in this field. Summary of the Invention

[0007] To address the aforementioned deficiencies of existing hot forging die steels, the present invention aims to provide a nano-high purity, high thermal stability hot forging die steel, TGE28. Through narrow-range internal composition control design, ultra-pure smelting control, and in-situ strengthening with nano-alloys, the purity, quenching and tempering temperature stability, and thermal fatigue performance of the material are simultaneously improved.

[0008] Another objective of this invention is to provide a method for preparing the aforementioned mold steel, which achieves uniform dispersion and high recovery rate of the nanophase in the steel matrix through a wire feeding addition process at the end of refining, thereby stably exerting the nano-reinforcing effect.

[0009] The objective of this invention can be achieved through the following technical solutions: This invention provides a method for preparing nano-high purity, high thermal stability hot forging die steel. The die steel has the following composition by mass percentage: C 0.38%–0.40%, Si 0.20%–0.40%, Mn 0.30%–0.50%, P≤0.015%, S≤0.002%, Cr 4.80%–5.20%, Mo 2.90%–3.20%, V 0.40%–0.60%, Ni≤0.25%, Cu≤0.10%. A 6m nano-alloy cored wire is added to each ton of molten steel using a wire feeding method, with the balance being Fe and unavoidable impurities. The method sequentially includes the following steps: medium-frequency induction melting, AOD refining, LF ladle refining, VD vacuum refining, protective casting to prepare consumable electrodes, electroslag remelting, forging, precision forging, and heat treatment.

[0010] The failure of hot forging die steel is essentially the result of the combined effects of high-temperature softening and thermal fatigue, and its performance is determined by the matrix composition, inclusion level, and microstructure uniformity. This invention employs a compositional system of "medium-carbon matrix, high-Mo secondary hardening, low-Si and low-Mn pure design, ultra-strict P, S, and O control, and in-situ strengthening with nano-alloys," enabling the matrix to maintain sufficient secondary hardening capacity during long-term service at 500-650℃. The process utilizes an ultra-pure flow of "medium-frequency induction, AOD, LF, VD, and electroslag remelting," with the nano-alloys added at the end of the LF stage using a cored wire feeding method. This avoids the burning and agglomeration of the nanophase during high-temperature oxidation, and further disperses the nanophases through subsequent VD soft stirring and electroslag remelting, thereby stably obtaining a high-purity, fine-grained, and dispersed microstructure on an industrial scale.

[0011] Furthermore, the medium-frequency induction melting uses low-phosphorus pure raw materials free of harmful elements, and no Mo, Cr, or V alloys are added in the medium-frequency furnace. The tapping temperature is ≥1630℃, and the tapping steel has C≥1.0% and P≤0.020%. The AOD refining adopts a double-slag method with argon protection throughout the process. After dephosphorization of the primary slag at 1520~1570℃, the oxide slag is discharged. High-carbon ferrochrome is added and the temperature is raised to 1680~1730℃ to decarburize to C≤0.20%. Then, Mo, Si, and Al are added for reduction, with a reduction strength of 1.1~1.3. The tapping temperature is ≥1600℃, and the tapping steel has Si≤0.10%.

[0012] Phosphorus (P) segregates at grain boundaries in steel, inducing temper brittleness and promoting intergranular propagation of thermal fatigue cracks. It must be efficiently removed under oxidizing conditions. However, oxides of alloying elements such as Cr, Mo, and V have high stability; if added before dephosphorization, they significantly increase the activity of oxides in the slag, inhibit the dephosphorization reaction, and cause alloy burn-off. This invention postpones alloying until after dephosphorization slag discharge. First, high-basicity, high-volume, high-FeO primary slag is used for low-temperature dephosphorization at 1520-1570℃. Then, oxygen blowing decarburization is performed at 1680-1730℃ using the strong stirring characteristics of AOD (Alternating Current Oxygen). After decarburization, Si and Al are used to strongly reduce Cr and Mo oxides in the slag, controlling the reduction intensity to 1.1-1.3, resulting in high alloy recovery and low oxygen levels in the molten steel. The Si content in the tapped steel is controlled below 0.10%, creating low-silicon conditions for subsequent deep aluminum deoxidation in LF (Leakage Processing), avoiding the decrease in thermal conductivity and temper brittleness caused by high Si content.

[0013] Furthermore, in the LF ladle refining process, after the molten steel enters the furnace, argon gas is introduced and the argon gas flow rate is adjusted to a state of slight fluctuation in the molten steel. In the early stage, aluminum wire is fed in at once to control the Al content in the steel to 0.08%~0.12% for sufficient deoxidation. During the smelting process, Si and C reducing agents are added in small amounts and multiple times for diffusion deoxidation. The white slag is maintained for ≥30min. In the later stage, the alloy composition is adjusted according to internal control and SiO2 is added to adjust the slag basicity to 2.4~2.6. 10~20min before VD, the nano-alloy cored wire is fed into the molten steel at a speed of 200m / min. Before tapping, the O content in the steel is controlled to ≤20ppm. Before the VD vacuum refining, the surface slag is removed and the free space of the ladle is ensured. During vacuum treatment, the vacuum degree is ≤30Pa, the high vacuum holding time is ≥10min, and the soft stirring time after breaking the vacuum is ≥20min.

[0014] Oxygen in steel mainly exists in the form of hard and brittle oxides such as Al2O3 and SiO2, which are natural sources of thermal fatigue crack initiation. In the early stages of LF (Furnace Drying), strong aluminum deoxidation reduces dissolved oxygen to extremely low levels, followed by continuous removal of inclusions through white slag diffusion deoxidation. The slag basicity is adjusted from the high value required for high basicity desulfurization to 2.4-2.6, improving slag fluidity and inclusion adsorption capacity while ensuring desulfurization capability. The nano-alloy is fed in 10-20 minutes before VD (Vacuum Deoxidation) because: at this time, the molten steel is already deeply deoxidized and has a low sulfur content, reducing the probability of nano-ceramic particles being entrained by oxide inclusions or forming sulfide nuclei; immediately after feeding, VD vacuum treatment is initiated, where the turbulence and soft stirring of the molten steel under vacuum provide the kinetic conditions for uniform dispersion of the nano-phase, while subsequent calming allows large inclusion particles to float to the surface, and the nano-phase to be dispersed and retained in the matrix. A vacuum degree of ≤30Pa and a holding pressure of ≥10min can reduce hydrogen and nitrogen to a safe range and promote the flotation of large particle inclusions. Soft stirring of ≥20min can further purify the molten steel and provide a guarantee of O≤20ppm and N≤120ppm.

[0015] Furthermore, argon gas protection is used throughout the casting process, and the casting temperature is the liquidus temperature of steel +45~50℃. After casting, the red-hot steel is annealed to obtain consumable electrode billets. The electroslag remelting adopts bipolar remelting. Before remelting, the slag and electrode billets are baked at a temperature of 200~500℃. The remelting slag system is a CaF2 to Al2O3 mass ratio of 80:20, and argon gas protection is used throughout the remelting process. The electrical parameters for the 1.3t ingot are 5500~7000A current and 63~65V voltage, and the electrical parameters for the 3.5t ingot are 7800~10000A current and 67~70V voltage. The slag pool temperature is stabilized at 1520~1560℃. The melting rate is controlled at 4.0~6.0kg / min, and the melting pool depth is stabilized at 40~60mm. The cooling water flow rate of the crystallizer is 15~20m³ / min. 3 / h, the inlet water temperature is not higher than the ambient temperature by 10℃, and the temperature difference between the inlet and outlet water is ≤10℃; during the final stage of remelting, the current is gradually reduced to compensate for shrinkage, and the current reduction process is set with a slope time of 20~30s. The total compensation time for 1.3t ingots is ≥30min, and the total compensation time for 3.5t ingots is ≥55min. During the final stage of compensation, the current is reduced to 20%~30% of the initial current.

[0016] To protect the casting process from secondary oxidation and nitrogen absorption of the molten steel, the casting superheat is controlled at 45-50°C above the liquidus line. This ensures the steel's filling capacity while preventing excessive superheat from leading to coarse, porous, and segregated solidification. Electroslag remelting is the key process for achieving ultrapure and dense microstructure in this invention: the CaF2-Al2O3 octahedral slag system has a well-matched melting point, viscosity, and conductivity, effectively adsorbing sulfide and oxide inclusions while controlling the morphology of carbide precipitation. Matching the current and voltage according to the ingot shape stabilizes the slag pool temperature at 1520-1560°C, and with a constant melting rate of 4.0-6.0 kg / min, the molten pool depth is stabilized at a shallow and flat shape of 40-60 mm, which is conducive to bottom-up directional solidification and inhibits element segregation and carbide network precipitation. Strong cooling of the crystallizer and control of the temperature difference between the inlet and outlet water ensure uniform solidification shell and prevent cracking. During the feeding stage, the current is reduced by ramping and the feeding time is extended, which can gradually fill the shrinkage at the top of the steel ingot, eliminate shrinkage cavities and porosity, and obtain a dense consumable ingot.

[0017] Furthermore, the forging process employs segmented heating: the steel ingot is first held at 650~700℃ for 2 hours, then heated to 800~850℃ and held for 1.5~2 hours, and finally heated to 1250~1270℃ and held for 16~18 hours; during forging, the upsetting ratio is ≥6, the drawing feed is ≤2 / 3 of the forging width, the single-pass reduction in rapid forging is controlled at 50~80mm, the final forging temperature is ≥850℃, and the billet is opened to the target size; after opening, homogenization annealing is performed: held at 780~800℃ for 8~10 hours, and then cooled to below 500℃ at a rate of ≤20℃ / h before being removed from the furnace.

[0018] Electroslag ingots still contain microscopic dendritic segregation and primary carbides of alloying elements such as Cr, Mo, and V. Segmented heating is employed, with preheating at 650-700℃ and 800-850℃ to reduce thermal stress and gradually transform the microstructure. A high-temperature hold at 1250-1270℃ for 16-18 hours further dissolves coarse primary carbides and ensures complete and uniform diffusion of elements. This temperature is below the solidus line and within the austenitic region, preventing overheating. A high upsetting ratio (≥6) combined with a fast forging process using small feed and large reduction ensures thorough three-dimensional breaking down of cast dendrites and coarse carbides, refining and uniformly distributing them along the deformation direction. A final forging temperature ≥850℃ prevents mixed crystals and cracks during forging in the two-phase region. Post-forging homogenization annealing at 780-800℃ further eliminates residual dendritic segregation and forging stress, while slow cooling prevents hydrogen-induced cracking and hard phase cracking.

[0019] Furthermore, precision forging is performed after holding at 1160~1180℃ using a ring furnace; after forging, it is quenched in water at above 850℃ for 3~5 minutes, followed by the following sequential processes: stress-relief annealing at 860℃ for 10 hours; ultra-fine treatment by water cooling after heating to 1010~1030℃ for 2~3 hours; austenitization by heating to 860~880℃ for 6~8 hours, followed by spheroidization by cooling to 730℃ at 10~15℃ / h and holding for 20 hours; and spheroidizing annealing by furnace cooling to below 500℃, with the hardness controlled at 180~220HBW after spheroidization; the heat treatment steps are: oil quenching at 1010~1070℃ for 30 minutes, followed by air tempering at 600~650℃ for 2 hours.

[0020] The precision forging temperature is selected at 1160~1180℃, which is in the full austenitic region and lower than the billet opening temperature. This allows for further grain refinement and improved dimensional accuracy under relatively low deformation resistance. Post-forging water quenching utilizes the high-temperature residual heat for rapid cooling to obtain a fine-grained structure, followed by stress-relief annealing to eliminate quenching stress. Ultra-refining treatment at 1010~1030℃, through re-austenitization and rapid cooling, further refines the original austenite grains and undissolved carbides, providing more nucleation sites for subsequent spheroidization. Spheroidizing annealing employs an isothermal spheroidization process involving slow cooling to 730℃ after austenitization and prolonged holding. This spheroidizes lamellar carbides into fine, uniform granular carbides, reducing the hardness to 180~220 HBW and ensuring good machinability. The final quenching temperature was selected at 1010~1070℃. As can be seen from the CCT curve, this temperature is in the full austenitic region and is higher than Ac3=914℃. The oil cooling rate can exceed the critical cooling rate of 0.4℃ / s to obtain full martensite. During tempering at 600~650℃, the high Mo and V content promotes the dispersion precipitation of M6C and M(C,N) nano carbides, resulting in significant secondary hardening. At the same time, it avoids the temperature range in which a large amount of Laves phase precipitates, ensuring a balance between strength and toughness.

[0021] Furthermore, the present invention also provides a nano-high purity, high thermal stability hot forging die steel prepared by the above method, with the following composition by mass percentage: C 0.38%–0.40%, Si 0.20%–0.40%, Mn 0.30%–0.50%, P≤0.015%, S≤0.002%, Cr 4.80%–5.20%, Mo 2.90%–3.20%, V 0.40%–0.60%, Ni≤0.25%, Cu≤0.10%. A 6m nano-alloy cored wire is added per ton of molten steel using a wire feeding method, with the balance being Fe and unavoidable impurities. The gas content of the finished product is controlled to be O≤20ppm and N≤120ppm.

[0022] The working principles of each element are as follows: C is the most basic strengthening element. On the one hand, it strengthens the matrix through solid solution and ensures hardness after quenching. On the other hand, it forms alloy carbides with Cr, Mo, and V to provide secondary hardening. In this invention, a narrow range of medium carbon (0.38%–0.40%) is used to balance hardness and toughness. Si assists in deoxidation and solid solution strengthening, but excessive Si reduces thermal conductivity, exacerbates temper brittleness, and promotes graphitization. In this invention, a low Si (0.20%–0.40%) is used to improve thermal conductivity, thereby accelerating mold heat dissipation, reducing thermal stress from temperature differences during hot and cold cycles, and alleviating thermal fatigue cracking. Mn assists in deoxidation and desulfurization and improves hardenability, but excessive Mn increases temper brittleness. In this invention, a low Mn (0.30%–0.50%) is used in combination with high Mo to offset temper brittleness. Cr is used at 4.80%–5.20% to ensure hardenability in large cross-sections and to form Cr7C3 and Cr... 23 C6 provides wear resistance and oxidation resistance. Mo, at 2.90%–3.20%, is nearly twice that of conventional H13 (approximately 1.1%–1.75%), forming fine spherical M6C and Mo2C that are extremely difficult to coarsen at high temperatures. This is the core source of resistance to temper softening and high-temperature strength, and it also suppresses temper brittleness caused by P and Mn. V, at 0.40%–0.60%, generates high-temperature stable VC, V4C3, and V(C,N), refining the original austenite grains, pinning dislocations and grain boundaries, and synergistically strengthening secondary hardening with Mo. P≤0.015% and S≤0.002% are controlled for ultra-low harmful elements, reducing the initiation of thermal fatigue cracks from both grain boundaries and inclusions. Cu≤0.10% prevents copper embrittlement and surface cracks. O≤20ppm and N≤120ppm are achieved through VD vacuum degassing, which reduces oxide inclusions while retaining the beneficial effect of a small amount of N in fine grains.

[0023] Furthermore, the nano-alloy cored wire contains TiCN and TiBN nano-ceramic particles; these nano-ceramic particles act as heterogeneous nucleation cores during the solidification of molten steel and subsequent hot working process, refining the as-cast microstructure and original austenite grains, pinning grain boundaries to inhibit high-temperature grain growth, and simultaneously hindering dislocation movement and delaying high-temperature recovery softening through the orovan dispersion strengthening mechanism.

[0024] TiCN and TiBN are both high-melting-point, high-hardness ceramic phases with low mismatch with austenite lattice. They can serve as effective heterogeneous nucleation substrates in the early stages of steel solidification, increasing the nucleation rate and refining the as-cast grains. During subsequent forging and quenching heating, these nanophases do not dissolve back, hindering austenite grain boundary migration through the Zener pinning mechanism, inhibiting grain coarsening, and allowing the steel to maintain a fine grain even at higher quenching temperatures. During cooling and tempering, the dispersed nanoparticles interact with dislocations, producing significant dispersion strengthening according to the Orovan mechanism, while simultaneously delaying dislocation climb and subgrain growth, thus postponing high-temperature recovery and recrystallization, and improving tempering stability. The nanophases can also serve as preferential nucleation sites for carbide precipitation, optimizing the morphology of M6C and M(C,N) precipitation, making them finer and more dispersed, reducing local stress concentration, improving thermal fatigue resistance, and enhancing the continuity of the surface oxide film and reducing high-temperature oxidation damage.

[0025] Furthermore, the mold steel, after being oil-quenched at 1010~1070℃ for 30 minutes and air-cooled for 2 hours, has a room temperature hardness of 56~60 HRC and a tempered hardness of 46~50 HRC; after standard heat treatment, its V-notch impact toughness is 320~350 J / cm². 2 .

[0026] The CCT curve calculated by JMatPro shows that the critical quenching rate of the steel of this invention is only about 0.4℃ / s, significantly lower than the approximately 1.0℃ / s of conventional H13. Oil quenching is sufficient to obtain full martensite in the core of large sections, and the quenching hardness remains stable at a high level. The 5% Cr and high Mo content strongly stabilize the supercooled austenite, significantly extending the incubation period for pearlite and bainite transformation. Normal quenching is less likely to produce bainite structures, avoiding the problem of poor high-temperature tempering stability of bainite. During tempering at 600~650℃, the high Mo, high V system precipitates M6C and M(C,N) nanocarbides, resulting in secondary hardening and maintaining the tempering hardness at 46~50 HRC. Fine-grained reinforcement (nanophase pinning to V(C,N)), fine and uniform carbides after spheroidizing annealing, and a low-inclusion matrix due to ultra-low S, P, and O content collectively ensure high impact toughness. The low-Si and low-Mn design reduces temper brittleness and grain boundary embrittlement, stabilizing the V-shaped impact toughness at 320~350 J / cm. 2 .

[0027] Furthermore, after the mold steel is subjected to long-term high-temperature tempering at 600~650℃ for 8~10h, the hardness retention rate is ≥85%; after self-constrained thermal fatigue test, under the same number of cold and hot cycles, the number, length and width of its surface thermal fatigue cracks are lower than those of H13 steel treated with the same process.

[0028] The resistance to tempering depends on the coarsening rate of the precipitated phase at high temperatures. The high-Mo M6C and Mo2C carbides formed in this invention exhibit extremely low Ostwald ripening rates at 600-650℃, with even higher V(C,N) high-temperature stability. Nano-TiCN / TiBN further pins dislocations and subgrain boundaries, delaying recovery softening. Therefore, the hardness decreases slowly and retains a high rate after prolonged tempering. Regarding thermal fatigue, cracks initiate at surface oxides / inclusions and propagate along grain boundaries and stress concentration zones: ultra-low S and O minimize MnS and oxide inclusions; low Si improves thermal conductivity and surface oxide film continuity; high Mo enhances high-temperature strength and creep resistance; and nano-phases refine grains and hinder crack propagation. The synergistic effect of these multiple factors delays crack initiation time and reduces propagation rate under thermal cycling, thus resulting in superior thermal fatigue resistance compared to H13.

[0029] The beneficial effects of this invention are: (1) This invention adopts a composition system of C 0.38%~0.40%, Si 0.20%~0.40%, Mn 0.30%~0.50%, P≤0.015%, S≤0.002%, Cr 4.80%~5.20%, Mo 2.90%~3.20%, V 0.40%~0.60%, Ni≤0.25%, Cu≤0.10%, and feeds 6m of nano-alloy cored wire containing TiCN and TiBN per ton of molten steel, with the balance being Fe and impurities, O≤20ppm and N≤120ppm. Through the synergistic design of "medium carbon, high Mo, low Si and low Mn, ultra-strict P / S / O, and nano-strengthening", combined with medium frequency melting, AOD double slag refining, LF deep aluminum deoxidation and final wire feeding, VD vacuum degassing, argon-protected casting, CaF2-A The ultra-pure preparation process of L2O3 slag-based electroslag remelting, large-scale segmented forging, and isothermal spheroidizing annealing significantly reduces harmful inclusions such as sulfides and oxides in steel, effectively controlling grain boundary weakening and thermal fatigue crack initiation sources. The M6C and M(C,N) nanocarbides formed by high Mo and V exhibit low coarsening rates and strong secondary hardening capabilities during long-term service at 500-650℃. The nano-ceramic phase, through heterogeneous nucleation to refine grains, Zener pinning of grain boundaries, and Orovan strengthening, hinders dislocation movement and delays high-temperature recovery softening. Therefore, this invention can obtain a fully martensitic structure and a uniformly fine spheroidized pre-structure within a wide quenching and tempering temperature window, with a good match between hardness, impact toughness, and tempering resistance. The hardenability is sufficient to support complete hardening of the core of large-tonnage, large-section hot forging dies, and the forging and heat treatment processes have high tolerance and good batch stability.

[0030] (2) The preparation method of this invention is compatible with existing special steel production lines. The nano-alloy is added at the end of LF refining by cored wire feeding and further dispersed by VD soft stirring and electroslag remelting, avoiding the burning and agglomeration of the nano-phase during high-temperature oxidation. The recovery rate and dispersion uniformity are high, and industrial mass production can be achieved without the need for new large-scale equipment. In electroslag remelting, the electrical parameters are matched according to the ingot shape, and the shallow flat molten pool and slope feeding are controlled by constant melting rate, which further ensures that the steel ingot is dense, uniform in composition, and has few shrinkage cavities. In summary, this invention improves the purity of composition, the stability of high-temperature structure and the resistance to thermal fatigue at the same time. It can effectively alleviate the common problems of high-temperature collapse, thermal fatigue cracking and early cracking of hot forging dies, and is suitable for large-tonnage precision hot forging dies. Attached Figure Description

[0031] The invention will now be further described with reference to the accompanying drawings.

[0032] Figure 1 The thermodynamic equilibrium phase diagram and precipitate simulation results of the TGE28 mold steel of this invention; Figure 2 The continuous cooling transformation (CCT) curve of the TGE28 mold steel of this invention; Figure 3 The microstructure of the TGE28 mold steel of this invention after quenching (left) and tempering (right) at 1010℃ is shown. Figure 4 The microstructure of the TGE28 mold steel of this invention at 1030℃ is shown in the left (quenched) and right (tempered) grain sizes. Figure 5 The microstructure of the TGE28 mold steel of this invention after quenching (left) and tempering (right) at 1050℃ is shown. Figure 6 The microstructure of the TGE28 mold steel of this invention after quenching (left) and tempering (right) at 1070℃ is shown. Figure 7 The hardness change curves of the TGE28 mold steel of this invention after long-term tempering at 600℃ and 650℃ at different quenching temperatures are shown. Figure 8 The microstructure of the TGE28 mold steel after spheroidizing annealing in this invention (50x, 100x and 500x magnification). Figure 9 This is a schematic diagram of the dimensions of the thermal fatigue test specimen of the present invention; Figure 10 The present invention relates to TGE28 mold steel after 200 and 500 cycles and the crack initiation morphology; Figure 11 The present invention describes the H13 steel after 200 and 500 cycles and the crack initiation morphology. Figure 12 This is the crack source energy spectrum of the TGE28 mold steel of the present invention. Detailed Implementation

[0033] The specific embodiments of the present invention will be described in detail below, but it should be understood that the scope of protection of the present invention is not limited to the specific embodiments.

[0034] Example 1 This embodiment prepares nano-high purity, high thermal stability hot forging die steel TGE28. The chemical composition of the finished product by mass percentage is: C 0.39%, Si 0.32%, Mn 0.41%, P 0.012%, S 0.0018%, Cr 5.02%, Mo 3.05%, V 0.51%, Ni 0.18%, Cu 0.07%. The nano-alloy cored wire feed rate is 6m / t of molten steel, with the balance being Fe and unavoidable impurities. The gas content of the finished product is O 18ppm and N 112ppm.

[0035] Preparation steps: S1 Medium-frequency induction melting: Low-phosphorus, pure scrap steel free of harmful elements is used as raw material. No Mo, Cr, or V alloys are added to the medium-frequency furnace to avoid alloy loss during subsequent dephosphorization processes. After melting, the steel is tapped at a temperature of 1645℃. The steel composition is: C 1.12%, P 0.018%, Ni / Cu ≤ 0.20%, with the balance being Fe and unavoidable impurities.

[0036] S2 AOD refining: The smelting process employs a dual-slag method with argon protection throughout. The primary slag is dephosphorized at 1550℃ until P ≤ 0.010%, after which the oxide slag is discharged. The slag is then reformed and high-carbon ferrochrome is added. The temperature is raised to 1700℃ for decarburization to C 0.18%, followed by the addition of Si and Al reducing agents for reduction, with a reduction strength of 1.2. After reduction, the slag is discharged and the steel is tapped at 1620℃ with a Si content of 0.08%.

[0037] S3 LF ladle refining: After the molten steel enters the LF furnace, argon gas is introduced for slight stirring. In the early stage, aluminum wire is fed to control the Al content to 0.10% for sufficient deoxidation, and the white slag state is maintained for 35 minutes. In the later stage, the alloy is adjusted to the internal control composition, and SiO2 is added to adjust the slag basicity to 2.5. 15 minutes before going to the VD furnace, nano-alloy cored wire is fed at a speed of 200 m / min, with a feeding rate of 6 m / t of molten steel. The oxygen content of the molten steel is controlled to 17 ppm before tapping.

[0038] S4 VD Vacuum Refining: After removing the surface slag, the molten steel is put into the VD furnace, and the vacuum degree is controlled at 25Pa. The high vacuum pressure is maintained for 12 minutes. After breaking the vacuum, argon gas is used for soft stirring for 25 minutes to promote the floating and adsorption of inclusions and improve the purity of the molten steel.

[0039] S5 Protective Casting: Argon gas is used for the entire casting process. The casting temperature is the liquidus temperature +48℃. After casting, the red steel is annealed in the furnace to obtain a 3.5t consumable electrode billet.

[0040] S6 Electroslag Remelting: Employs a bipolar remelting process, with the slag material thoroughly baked before remelting and the electrode blank baked at 350℃; uses a CaF2:Al2O3 = 80:20 slag system, with argon protection throughout the process; controls the current at 8500A, voltage at 68V, melting rate at 5.0kg / min, and melt pool depth at approximately 50mm; the crystallizer cooling water flow rate is 18m³ / min. 3 / h, inlet and outlet water temperature difference 8℃; during the final stage of remelting, the current is gradually reduced to compensate for shrinkage, the compensation time is 50min, the current is reduced to 25% of the initial value at the end, and the current reduction ramp time is 25s.

[0041] S7 Forging Billet Opening: The steel ingot is heated in stages: 680℃ for 2 hours, 820℃ for 1.8 hours, and 1260℃ for 17 hours; a 2000t high-speed forging mill is used for billet opening, with an upsetting ratio of 6, a drawing feed of 1 / 2 of the forging width, a single-pass reduction of 65mm, and a final forging temperature of 880℃, opening the billet to a cross-section of 110mm×110mm; after billet opening, it is homogenized by holding at 790℃ for 9 hours, and then cooled in the furnace at a cooling rate of 18℃ / h to below 480℃ before being unloaded.

[0042] S8 Precision forging and finished product heat treatment: The billet is precision forged to the finished size after being held at 1170℃; after forging, it is quenched in water at 860℃ or above for 4 minutes, followed by stress relief annealing at 860℃ for 10 hours and water cooling ultrafine treatment at 1020℃ for 2.5 hours; finally, the finished product is spheroidized annealed: after austenitization at 870℃ for 7 hours, it is cooled to 730℃ at 12℃ / h and held for 20 hours to spheroidize, and then furnace cooled to below 480℃ before being taken out of the furnace. The spheroidized hardness is about 200HBW.

[0043] The final heat treatment process for the finished product is as follows: oil quenching at 1030℃ for 30 minutes, followed by air tempering at 605℃ for 2 hours, for a total of 3 tempering cycles.

[0044] The equilibrium phase calculations for the composition of this invention were performed using JMatPro V11.0 thermodynamic software (see...). Figure 1 ) and CCT curve simulation (see Figure 2 ).

[0045] (1) Equilibrium Phase Diagram: The steel of this invention begins to precipitate primary austenite at approximately 1420℃, and the liquid phase completely disappears at approximately 1380℃, entering the all-austenite single-phase region. The stable range of the all-austenite is approximately 560℃, providing a wide heating window for forging and quenching, and a high process tolerance. The forging heating temperature is controlled at 1250~1270℃, and the quenching temperature is controlled at 1010~1070℃, both falling within the pure austenite region, with no proprecipitated ferrite, ensuring complete quenching. When cooled to approximately 1350℃, V(C,N) vanadium carbonitrides precipitate first, exhibiting extremely strong high-temperature stability. They do not re-dissolve during forging / quenching heating, pinning austenite grain boundaries and inhibiting grain coarsening. 820℃ is the equilibrium phase transformation point, where austenite decomposes in large quantities, and M... 23 Secondary carbides such as C6 and M6C precipitate explosively; the Laves phase (the brittle intermetallic phase of Mo) precipitates stably below 750℃, therefore the tempering temperature should not be higher than 650℃ for extended periods, and the service temperature should be avoided from stabilizing above 700℃. In the 500~700℃ range, M... 23 C6 provides basic wear resistance and oxidation resistance, while M6C molybdenum carbide is fundamental to resisting temper softening. V(C,N) nanophase pinning dislocations improves the resistance to thermal fatigue crack propagation. These three elements work synergistically to achieve a balance between strength, toughness, and stability. Compared to H13, the steel of this invention has a higher upper limit of austenite stability, a larger amount of M6C precipitation, a higher V(C,N) precipitation temperature, and very few impurity phases such as MnS / M3P, resulting in superior processability, tempering stability, and fatigue resistance.

[0046] (2) CCT curve: The pearlite transformation nose point is located in the extremely slow cooling rate region with a cooling rate ≤ 0.02℃ / s, and the incubation period is extremely long. Lamellar pearlite hardly appears during conventional heat treatment and post-forging cooling; the bainite transformation region is 353~441℃, corresponding to a cooling rate of 0.02~0.4℃ / s. This critical cooling rate range should be avoided during quenching to prevent bainite from causing insufficient hardness at high temperatures; the martensite initiation temperature M s At approximately 380℃, a cooling rate ≥0.4℃ / s is sufficient to obtain full martensite. The critical cooling rate is much lower than that of H13, allowing for complete hardening of the core in large cross-sections without a soft core layer. Ac3 = 914℃, and the quenching temperature of 1010~1070℃ is much higher than Ac3, ensuring complete austenitization. Oil cooling is sufficient to meet the cooling rate requirements, reducing the risk of quenching cracks. Post-forging softening annealing requires furnace cooling (cooling rate ≤0.02℃ / s) to fully precipitate spheroidized structures and reduce hardness for easier machining.

[0047] Performance test results: (1) Quenching and tempering hardness: The samples were oil quenched at 1010, 1030, 1050 and 1070℃ for 30 min respectively, and then tempered once at 590℃ for 2 h and twice at 620℃ for 2 h. The results are shown in Table 1. The electron microscope images of the quenched microstructure (left) and tempered grain size (right) at different temperatures are shown in Table 1. Figure 3 (1010℃) Figure 4 (1030℃) Figure 5 (1050℃) and Figure 6 As shown in (1070℃), the quenching hardness is stable at 57~59HRC within the quenching range of 1010~1070℃. After secondary tempering, it still maintains a high hardness, and the hardness retention rate of tempering increases with the increase of quenching temperature.

[0048] Table 1. Quenching and tempering hardness of TGE28 at different quenching temperatures

[0049] (2) Impact toughness: After quenching at 1030℃ and tempering at 605℃ for 2 hours for 3 times, V-notch impact test specimens were taken transversely at the middle and head of the steel. The results are shown in Table 2. The average impact energy at the middle was 344J, at the head was 320J, and the overall energy ranged from 320 to 350J / cm. 2 The interval has a small difference between the head and the middle, and the organization is uniform.

[0050] Table 2 Impact Performance Test Results of TGE28

[0051] (3) Tempering resistance stability: After standard heat treatment, the hardness change curves after long-term holding at 600℃ and 650℃ are shown in the figure. Figure 7 It can be seen that the samples quenched at 1030℃ and 1050℃ still maintain high hardness after tempering at 600℃ for 50 hours, and have a high hardness retention rate after tempering at 650℃ for 10 hours, indicating good resistance to tempering softening.

[0052] (4) Microstructure observation: The metallographic structures at different magnifications after spheroidizing annealing are shown in the figure. Figure 8 (50x, 100x and 500x respectively), the microstructure is uniform and fine, the carbides are well spheroidized, and there is no obvious segregation or coarse carbides, which provides a good pre-structure for subsequent heat treatment.

[0053] Example 2 The composition of this embodiment by mass percentage is as follows: C 0.38%, Si 0.22%, Mn 0.32%, P 0.013%, S 0.0015%, Cr 4.85%, Mo 2.95%, V 0.42%, Ni 0.15%, Cu 0.06%, nano-alloy feed rate 6m / t, balance Fe and impurities; O 19ppm, N 108ppm.

[0054] The preparation process differs from that in Example 1 in that: the electroslag remelting is carried out on a 1.3t ingot with a current of 6000A and a voltage of 64V for 35 minutes; the final heat treatment is oil quenching at 1010℃ for 30 minutes and air tempering at 600℃ for 2 hours.

[0055] Properties: Quenched hardness 57.3 HRC, tempered hardness 48.5 HRC, V-notch impact toughness 348 J / cm² 2 The hardness retention rate after tempering at 600℃ for 10 hours is approximately 87%.

[0056] Example 3 The composition of this embodiment by mass percentage is as follows: C 0.39%, Si 0.30%, Mn 0.40%, P 0.011%, S 0.0017%, Cr 5.00%, Mo 3.08%, V 0.50%, Ni 0.20%, Cu 0.08%, with a nano-alloy feed rate of 6 m / t, and the balance being Fe and impurities; O 16 ppm, N 115 ppm.

[0057] The preparation process differs from that in Example 1 in that the final heat treatment is oil quenching at 1050℃ for 30 min and air tempering at 620℃ for 2 h.

[0058] Properties: Quenched hardness 57.4 HRC, first tempered hardness 55.1 HRC, second tempered hardness 47.9 HRC, V-notch impact toughness 335 J / cm² 2 The hardness retention rate after tempering at 620℃ for 10 hours is approximately 88%. This example covers the median composition and tempering temperature, and exhibits stable microstructure and properties.

[0059] Example 4 The composition of this embodiment by mass percentage is as follows: C 0.40%, Si 0.38%, Mn 0.48%, P 0.014%, S 0.0019%, Cr 5.18%, Mo 3.18%, V 0.58%, Ni 0.22%, Cu 0.09%, with a nano-alloy feed rate of 6 m / t, and the balance being Fe and impurities; O 20 ppm, N 118 ppm.

[0060] The preparation process differs from that in Example 1 in that the final heat treatment is oil quenching at 1070℃ for 30 min and air tempering at 650℃ for 2 h.

[0061] Properties: Quenched hardness 59.4 HRC, first tempered hardness 56.1 HRC, second tempered hardness 48.5 HRC, V-notch impact toughness 322 J / cm² 2 The hardness retention rate after tempering at 650℃ for 10 hours is approximately 85%.

[0062] Example 5 The composition of this embodiment is the same as that of Embodiment 3, the only difference being the process verification of the timing and speed of feeding the nano-alloy cored wire: 10 minutes before VD (Vacuum Dryer), feed at 200 m / min for 6 m / t (Scheme A); 20 minutes before VD (Vacuum Dryer), feed at 200 m / min for 6 m / t (Scheme B); 15 minutes before VD (Vacuum Dryer), feed at 180 m / min for 6 m / t (Scheme C). All three schemes can achieve uniform dispersion of the nano-phase, with finished product O ≤20 ppm, and quenched and tempered hardness and impact toughness comparable to Embodiment 3.

[0063] The results show that the nano-alloy can stably exert a strengthening effect within the window of 10-20 min before VD, about 200 m / min, and 6 m / t. If it is fed in at the beginning of LF (see Comparative Example 6), the nano-phase is easily entrained in the oxide slag and the recovery rate decreases.

[0064] Example 6 In this embodiment, 1.3t ingots and 3.5t ingots were used to prepare steel with the same composition (the composition is the same as in Example 3). The forging billet was a 110mm×110mm square billet. After the same spheroidizing annealing, quenching at 1030℃ and tempering at 605℃, samples were taken from the edge, half radius and center of the square billet cross section to test the hardness and impact toughness.

[0065] Results: The hardness difference between the edge and core of the 1.3t ingot was less than 1.5 HRC, and the impact energy fluctuation was less than 20 J; the hardness difference between the edge and core of the 3.5t ingot was less than 2 HRC, and the impact energy fluctuation was less than 25 J, while the core hardness still reached above 46 HRC and the impact energy was still above 320 J. This indicates that the steel of this invention has a low critical cooling rate and good hardenability. The electroslag remelting shallow flat molten pool and the high forging ratio forging ensure the uniform microstructure and properties of the core and edge of the large cross-section, meeting the requirements of large-tonnage dies.

[0066] Example 7 The composition of this embodiment is the same as that of Embodiment 1. After forging and blanking, different spheroidizing annealing processes are used: Scheme A (austenitizing at 870℃ for 7h, then cooling to 730℃ at 12℃ / h and holding for 20h), Scheme B (austenitizing at 860℃ for 8h, then cooling to 730℃ at 10℃ / h and holding for 20h), and Scheme C (austenitizing at 880℃ for 6h, then cooling to 730℃ at 15℃ / h and holding for 20h).

[0067] The hardness of all three spheroidized schemes was between 185 and 215 HBW. Metallographic observation showed that the carbides were fine and uniform granules, without obvious lamellar or network carbides. After subsequent quenching at 1030℃ and tempering at 605℃, the hardness and impact toughness were comparable to those of Example 1. This indicates that good spheroidized structure and machinability can be obtained within the spheroidizing process window of holding at 860-880℃ for 6-8 hours and cooling to 730℃ at a rate of 10-15℃ / h and holding for 20 hours.

[0068] Comparative Example 1 Traditional H13 hot work die steel was used, with the composition shown in Table 3. It underwent the same forging and spheroidizing annealing process as in Example 1, followed by standard heat treatment: oil quenching at 1030℃ and tempering at 605℃. A self-constrained thermal fatigue test was used to compare the thermal fatigue performance of TGE28 (Example 1) and H13. Specimen dimensions are shown in Table 3. Figure 8 The surface SEM morphology and crack source energy spectrum after cycling are shown in the figure. Figure 10-12 .

[0069] Table 3 Chemical composition (wt%) of samples from thermal fatigue test

[0070] Results: Under the same number of cycles, TGE28 had fewer surface cracks, and the cracks were finer and shallower. The oxidation and inclusions at the crack initiation sites of TGE28 were also less severe. Figure 10 The surface of H13 showed obvious network cracks and main crack propagation. Significant oxidation and inclusions were visible at the crack initiation points of H13. Figure 11 ); Energy spectrum ( Figure 12 The crack initiation site shows the presence of oxide / oxide layers such as O, Al, and Si. This demonstrates that the low-Si, high-Mo composition, ultra-pure design, and nano-reinforcement of this invention significantly improve thermal fatigue resistance.

[0071] Comparative Example 2 The composition and process of this comparative example are the same as those of Example 1, except that no nano-alloy cored wire is fed in at the end of the LF refining process, while the rest of the process remains unchanged.

[0072] Results: The composition and purity of the finished product were comparable to those of Example 1, but the original austenite grains were slightly coarser. The quenched and tempered hardness was close to that of Example 1 (approximately 57 HRC / 47 HRC), and the V-shaped impact toughness decreased to approximately 280 J / cm. 2 The hardness retention rate after tempering at 650℃ for 10 hours decreased to approximately 78%, and thermal fatigue crack initiation occurred earlier. This indicates that the nano-alloys make substantial contributions to toughness and tempering resistance through grain refinement strengthening, grain boundary pinning, and Orovan strengthening.

[0073] Comparative Example 3 The composition of this comparative example is the same as that of Example 1, but the Mo content is reduced to 1.50% (close to the H13 level), while the other elements and processes remain unchanged.

[0074] Results: The quenched hardness was approximately 56 HRC, but the secondary hardening peak decreased significantly after tempering above 600℃. The hardness retention rate after tempering at 650℃ for 10 hours was only about 70%, indicating insufficient resistance to high-temperature collapse. The impact toughness was approximately 300 J / cm. 2 This indicates that the Mo content is the core of resistance to tempering stability, and a Mo content of 2.90% to 3.20% is necessary to ensure the retention rate of high-temperature hardness.

[0075] Comparative Example 4 The composition and process of this comparative example are the same as those of Example 1, but conventional purity control is used: P about 0.025%, S about 0.015%, O about 35ppm, without VD high vacuum pressure holding and electroslag remelting purification.

[0076] Results: Metallographic and SEM observations revealed numerous MnS and oxide inclusions, and the V-shaped impact toughness decreased to approximately 220 J / cm². 2 In the early stages of the thermal fatigue test, network cracks originating from inclusions appeared, and the tempering stability also decreased. This indicates that ultrapure control with S≤0.002%, P≤0.015%, and O≤20ppm plays a decisive role in toughness and thermal fatigue life.

[0077] Comparative Example 5 The composition of this comparative example is the same as that of Example 1, but Si is increased to 1.00% (H13 level), while the other elements and processes remain unchanged.

[0078] Results: The hardness after quenching and tempering was comparable to that of Example 1, but the thermal conductivity of the steel decreased, the tendency for temper brittleness increased, and the V-shaped impact toughness decreased to approximately 260 J / cm. 2 During thermal fatigue cycles, the surface oxide film is easily peeled off, and cracks propagate faster. This indicates that low-Si design (0.20%–0.40%) is of substantial significance in improving thermal conductivity and alleviating thermal fatigue cracking.

[0079] Comparative Example 6 The composition of this comparative example is the same as that of Example 1, except that the nano-alloy cored wire is fed in at the early stage of LF (before aluminum deoxidation), while the rest of the process remains unchanged.

[0080] Results: In the initial stage of LF steel, the oxygen level was high and the slag volume was large. The nano-ceramic particles were easily carried and floated by the oxide slag, resulting in a significant decrease in yield. The nano-phase distribution in the billet was uneven, and local agglomeration occurred. The V-shaped impact toughness of the finished product decreased to about 290 J / cm. 2 The hardness retention rate after tempering at 650℃ for 10 hours was approximately 80%, which is lower than that of Example 1. This indicates that the nano-alloy should be fed into the molten steel 10-20 minutes before VD to ensure recovery rate and dispersion uniformity.

[0081] Comparative Example 7 The composition of this comparative example is the same as that of Example 1, except that the electroslag remelting process is omitted, and the electrode blank obtained by casting is directly homogenized and annealed and forged into finished products. The rest of the process remains unchanged.

[0082] Results: The steel contained a small amount of coarse oxides and sulfides, with significant central porosity and compositional segregation, and the V-shaped impact toughness decreased to approximately 270 J / cm. 2The hardness difference between the core and edge of the cross-section exceeds 3 HRC, resulting in a decrease in thermal fatigue life. This indicates that electroslag remelting plays an irreplaceable role in achieving ultrapure, dense microstructure and uniform large cross-sections.

[0083] Comparative Example 8 The composition of this comparative example is the same as that of Example 1, except that the upsetting ratio during forging is reduced to 3 (not reaching ≥6), and the drawing feed and reduction are correspondingly reduced, while the other heat treatment processes remain unchanged. Results: Metallographic observation shows incompletely broken coarse primary carbides and residual cast dendrites, resulting in an uneven microstructure; the V-shaped impact toughness is reduced to approximately 250 J / cm. 2 The transverse properties deteriorate significantly, and thermal fatigue cracks tend to propagate along coarse carbides. This indicates that a large upsetting ratio (≥6) combined with small feed and large reduction in rapid forging is essential for breaking down the as-cast microstructure and coarse carbides, and for ensuring isotropic toughness.

[0084] The above-disclosed embodiments are merely a few specific examples of the present invention. However, the embodiments of the present invention are not limited thereto, and any variations that can be conceived by those skilled in the art should fall within the protection scope of the present invention.

Claims

1. A method for preparing nano-high purity, high thermal stability hot forging die steel, characterized in that, The mold steel has the following composition by mass percentage: C 0.38%~0.40%, Si 0.20%~0.40%, Mn 0.30%~0.50%, P≤0.015%, S≤0.002%, Cr 4.80%~5.20%, Mo 2.90%~3.20%, V 0.40%~0.60%, Ni≤0.25%, Cu≤0.10%. 6m of nano-alloy cored wire is added per ton of molten steel using a wire feeding method, with the balance being Fe and unavoidable impurities. The method sequentially includes the following steps: medium-frequency induction melting, AOD refining, LF ladle refining, VD vacuum refining, protective casting to produce consumable electrodes, electroslag remelting, forging, precision forging, and heat treatment.

2. The method for preparing nano-high-purity, high-thermal-stability hot forging die steel according to claim 1, characterized in that, The medium-frequency induction melting uses low-phosphorus pure raw materials free of harmful elements. No Mo, Cr, or V alloys are added to the medium-frequency furnace. The tapping temperature is ≥1630℃, and the steel composition meets the requirements of C≥1.0%, P≤0.020%, and Ni / Cu≤0.25%. The AOD refining adopts a double-slag method with argon protection throughout the process. The primary slag is dephosphorized at 1520~1570℃. After dephosphorization, the oxide slag is removed from the mirror surface. High-carbon ferrochrome is added and the temperature is raised to 1680~1730℃ for decarburization to C≤0.20%. Then, Mo, Si, and Al alloys are added for reduction. The reduction intensity is controlled at 1.1~1.

3. After reduction, the slag is discharged and the temperature is raised for tapping. The tapping temperature is ≥1600℃, and the steel composition is ≤0.10%.

3. The method for preparing nano-high-purity, high-thermal-stability hot forging die steel according to claim 1, characterized in that, In the early stage of the LF ladle refining process, aluminum wire is fed in a single batch to control Al content to 0.08%~0.12% for sufficient deoxidation, and white slag is maintained for ≥30 minutes. In the later stage, SiO2 is added to adjust the basicity to 2.4~2.

6. 10~20 minutes before VD, 6m / t of nano-alloy cored wire is fed in at 200m / min. O ≤20ppm before tapping. The vacuum degree of VD vacuum refining is ≤30Pa, high vacuum pressure is maintained for ≥10 minutes, and soft stirring is performed for ≥20 minutes after breaking the vacuum.

4. The method for preparing nano-high-purity, high-thermal-stability hot forging die steel according to claim 1, characterized in that, The entire process of protective casting is under argon protection, and the casting temperature is liquidus +45~50℃; the electroslag remelting adopts bipolar remelting, the slag system is CaF2:Al2O3=80:20, the electrode billet is baked at 200~500℃, and the entire process is under argon protection. For 1.3t ingots, the current is 5500~7000A and the voltage is 63~65V; for 3.5t ingots, the current is 7800~10000A and the voltage is 67~70V; the slag pool temperature is 1520~1560℃; the melting rate is 4.0~6.0kg / min; the melt pool depth is 40~60mm; and the crystallizer water flow rate is 15~20m³ / min. 3 / h, inlet and outlet water temperature difference ≤10℃; the current at the end of the feeding period is reduced to 20~30% of the initial current, the feeding of 1.3t ingot is ≥30min, the feeding of 3.5t ingot is ≥55min, and the current reduction is set to a 20~30s slope.

5. The method for preparing nano-high-purity, high-thermal-stability hot forging die steel according to claim 1, characterized in that, The forging process employs segmented heating: holding at 650~700℃ for 2 hours, holding at 800~850℃ for 1.5~2 hours, and holding at 1250~1270℃ for 16~18 hours; upsetting ratio ≥6, drawing feed ≤2 / 3 of forging width, rapid forging with a single pass reduction of 50~80mm, and final forging temperature ≥850℃; after forging, homogenization annealing is performed at 780~800℃ for 8~10 hours, followed by cooling to below 500℃ at a rate of ≤20℃ / h before unloading from the furnace.

6. The method for preparing nano-high-purity, high-thermal-stability hot forging die steel according to claim 1, characterized in that, The precision forging is carried out after holding at 1160~1180℃; after forging, it is quenched in water at above 850℃ for 3~5 minutes, followed by stress relief annealing at 860℃ for 10 hours, then water-cooled for ultrafine processing at 1010~1030℃ for 2~3 hours, and finally austenitized at 860~880℃ for 6~8 hours, then cooled to 730℃ at 10~15℃ / h and held for 20 hours for spheroidization, and then furnace-cooled to below 500℃ for removal from the furnace, with a spheroidized hardness of 180~220HBW; the final heat treatment of the finished product is oil quenching at 1010~1070℃ for 30 minutes and air-cooled tempering at 600~650℃ for 2 hours.

7. A nano-high purity, high thermal stability hot forging die steel, characterized in that, Prepared by the preparation method according to any one of claims 1 to 6, the composition by mass percentage is: C 0.38% to 0.40%, Si 0.20% to 0.40%, Mn 0.30% to 0.50%, P≤0.015%, S≤0.002%, Cr 4.80% to 5.20%, Mo 2.90% to 3.20%, V 0.40% to 0.60%, Ni≤0.25%, Cu≤0.10%, 6m of nano-alloy cored wire is fed per ton of molten steel, the balance being Fe and unavoidable impurities, and the gas content is O≤20ppm and N≤120ppm.

8. The nano-high purity, high thermal stability hot forging die steel according to claim 7, characterized in that, The nano-alloy cored wire contains TiCN and TiBN nano-ceramic particles, which act as heterogeneous nucleation cores in steel to refine the as-cast microstructure and original austenite grains, pin grain boundaries to inhibit high-temperature grain growth, and hinder dislocation movement and delay high-temperature recovery softening through orovan dispersion strengthening.

9. The nano-high purity, high thermal stability hot forging die steel according to claim 7, characterized in that, The mold steel, after being oil-quenched at 1010~1070℃ for 30 minutes and air-cooled for tempering at 600~650℃ for 2 hours, has a room temperature hardness of 56~60HRC and a tempered hardness of 46~50HRC; after standard heat treatment, its V-notch impact toughness is 320~350J / cm². 2 .

10. The nano-high purity, high thermal stability hot forging die steel according to claim 7, characterized in that, The mold steel retains ≥85% of its hardness after being tempered at 600~650℃ for 8~10 hours; after self-constrained thermal fatigue testing, the number, length and width of surface cracks are lower than those of H13 steel produced by the same process under the same number of thermal cycles.

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