A 1700mpa grade ultra-high toughness and ultra-high strength steel and a preparation method thereof

Abstract

The application discloses a 1700MPa-grade super-high toughness super-high strength steel and a preparation method thereof, and belongs to the technical field of super-high strength steel, and solves the problems of high alloy element content, high cost and poor strength and toughness matching of the existing super-high strength steel.The 1700MPa-grade super-high toughness super-high strength steel comprises the following components in percentage by mass: C: 0.27% to 0.32%, Si: 1.30% to 1.60%, Mn: 0.50% to 0.80%, Ni: 0.80% to 1.20%, Cr: 3.20% to 3.60%, Mo: 0.31% to 0.60%, W: 0.40% to 0.80%, Nb: 0.001% to 0.009%, P: ≤0.005%, S: ≤0.001%, and the balance is iron and inevitable impurities.The application realizes the effect of preparing the super-high strength steel at low cost and high quality.

Description

Technical Field

[0001] This invention relates to the field of ultra-high strength steel technology, and in particular to a 1700MPa grade ultra-high toughness ultra-high strength steel and its preparation method. Background Technology

[0002] Ultra-high strength steel, due to its excellent specific strength and superior strength-toughness combination, has been widely used in key load-bearing components for many years, such as aircraft landing gear, engine shafts, and gears. In recent years, with the development of high-performance and low-cost equipment, the market has become highly sensitive to raw material costs. Traditional grades such as secondary hardening steel and martensitic aging steel, which rely on high Co / Ni alloy content and ultra-pure smelting processes to achieve ultra-high strength and high toughness, have faced limitations in their application. Against this backdrop, achieving a strength and toughness level comparable to high-alloy ultra-high strength steel by rationally designing smelting and hot-working process parameters within the composition system of medium- and low-alloy ultra-high strength steel is of great significance. Summary of the Invention

[0003] In view of the above, the present invention aims to provide a 1700MPa grade ultra-high toughness ultra-high strength steel and its preparation method, in order to solve at least one of the following problems: existing ultra-high strength steels have high alloy element content, high cost, and poor strength-toughness matching.

[0004] The objective of this invention is mainly achieved through the following technical solutions:

[0005] On one hand, the present invention provides a 1700MPa grade ultra-high toughness and ultra-high strength steel. The composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, includes: C: 0.27% to 0.32%, Si: 1.30% to 1.60%, Mn: 0.50% to 0.80%, Ni: 0.80% to 1.20%, Cr: 3.20% to 3.60%, Mo: 0.31% to 0.60%, W: 0.40% to 0.80%, Nb: 0.001% to 0.009%, P: ≤0.005%, S: ≤0.001%, with the balance being iron and unavoidable impurities.

[0006] Furthermore, in the composition of 1700MPa grade ultra-high toughness and ultra-high strength steel, Cr+Ni≥4.2% and Cr / Ni≥2.8, where Cr and Ni refer to the mass percentage of the corresponding elements.

[0007] Furthermore, in the composition of 1700MPa grade ultra-high toughness and ultra-high strength steel, W+Mo≥0.85% and W / Mo≥0.9, where W and Mo refer to the mass percentage of the corresponding elements.

[0008] Furthermore, the composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, includes: C: 0.27%–0.31%, Si: 1.32%–1.55%, Mn: 0.53%–0.75%, Ni: 0.90%–1.20%, Cr: 3.30%–3.58%, Mo: 0.40%–0.60%, W: 0.40%–0.70%, P: ≤0.005%, S: ≤0.001%, Nb: 0.005%–0.009%, with the balance being iron and unavoidable impurities.

[0009] Furthermore, the microstructure of the 1700MPa grade ultra-high toughness and ultra-high strength steel consists of lath martensite matrix + a small amount of thin film retained austenite and NbC particles, and finely dispersed ε-carbides.

[0010] Furthermore, the yield strength ratio of 1700MPa grade ultra-high toughness ultra-high strength steel is 0.77 to 0.80.

[0011] The present invention also provides a method for preparing the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, comprising: electric furnace melting; ladle refining; electrode rod casting; vacuum consumable remelting or electroslag remelting to obtain steel ingots; steel ingot annealing; steel ingot heating homogenization treatment; forging; and heat treatment.

[0012] Furthermore, in vacuum consumable remelting or electroslag remelting, the melting rate v during the steady-state melting stage is related to the ingot diameter D as follows: v = (0.0115 ~ 0.0125) × D - 3.

[0013] Furthermore, forging includes rapid forging, with a final forging deformation ratio ≥2 and a final forging temperature range of 780–950℃.

[0014] Furthermore, the heat treatment includes annealing, quenching, and tempering; the annealing holding temperature is 600-700℃, the quenching holding temperature is 900-950℃, and the tempering holding temperature is 200-300℃.

[0015] Compared with the prior art, the present invention can achieve at least one of the following beneficial effects:

[0016] a) In the 1700MPa grade ultra-high toughness and ultra-high strength steel of the present invention, the excellent performance of the steel is ensured by precisely controlling the types and contents of alloying elements such as Cr, Ni, Mo, W and Nb, as well as the synergistic quantitative relationship of different elements. The present invention does not contain precious elements such as Co, and the contents of elements such as Ni and Nb are very low, thus achieving the effect of preparing ultra-high strength steel with low raw material cost.

[0017] b) The preparation method of the 1700MPa grade ultra-high toughness and ultra-high strength steel of the present invention achieves stable production of large-size vacuum consumable ingots of Φ660mm and above by precisely controlling the key parameters such as the melting rate of vacuum self-consumable remelting or electroslag remelting steady-state melting and matching with the ingot shape. The steel ingot has uniform composition and light micro-segregation, avoiding the precipitation of large-size carbides caused by element agglomeration, which would damage toughness.

[0018] c) The preparation method of the present invention achieves the excellent properties of steel by precisely controlling the homogenization treatment, forging and other process steps and parameters, avoiding the precipitation of needle-like harmful phases during forging and thus avoiding damage to toughness, thereby ensuring the combination of ultra-high strength and ultra-high toughness.

[0019] d) The steel of this invention exhibits excellent properties and can meet application requirements. For example, its tensile strength is 1700 MPa or more (e.g., 1718–1780 MPa), its yield strength is 1350 MPa or more (e.g., 1351–1395 MPa), its yield strength ratio is between 0.77 and 0.80, its elongation A is 12% or more (e.g., 12.5%–15%), its reduction of area Z is 52% or more (e.g., 54%–61%), its impact energy KU2 is 80 J or more (e.g., 84–110 J), and its fracture toughness K... IC 150 MPa·m 1 / 2 Above (e.g., 167–190 MPa·m) 1 / 2 ).

[0020] Other features and advantages of the invention will be set forth in the description which follows, and will be apparent in part from the description, or may be learned by practicing the invention. The objects and other advantages of the invention may be realized and obtained by means of what is particularly pointed out in the written description and the accompanying drawings. Attached Figure Description

[0021] The accompanying drawings are for illustrative purposes only and are not intended to limit the invention. Throughout the drawings, the same reference numerals denote the same parts.

[0022] Figure 1 The microstructure (martensitic laths and thin austenite film) of Example 1;

[0023] Figure 2 The microstructure of Example 1 (mainly reinforcing phase ε-carbide);

[0024] Figure 3 The strain field at the crack tip (plastic deformation zone) during the fracture toughness test in Example 2 is shown.

[0025] Figure 4 The microstructure of Example 5 (nanoscale TiC can be seen);

[0026] Figure 5 The microstructure of Comparative Example 3 (containing large M6C particles);

[0027] Figure 6 For comparative example 4, the stress field at the crack tip (plastic deformation zone) during fracture toughness testing is shown.

[0028] Figure 7 The microstructure of Comparative Example 5 (containing needle-like AlN) is shown. Detailed Implementation

[0029] Preferred embodiments of the present invention will now be described in detail with reference to the accompanying drawings, which form part of the present invention and, together with the embodiments of the present invention, serve to illustrate the principles of the present invention.

[0030] Ultra-high strength steel is highly sensitive to impurities or second phases and the microcracks they induce. To improve material cleanliness and avoid large-sized precipitates, thereby enhancing toughness, high-alloy ultra-high strength steel is often smelted using the costly "double vacuum" (vacuum induction + vacuum arc remelting) process. Generally, medium- and low-alloy ultra-high strength steel is smelted using a lower-cost process of electric furnace + ladle refining + vacuum arc remelting / electroslag remelting to control overall costs. This requires, on the one hand, meticulous control of the refining process to obtain highly clean steel ingots, and on the other hand, optimized design of the hot working process to obtain a high-quality original forging structure.

[0031] This invention provides a 1700MPa grade ultra-high toughness and ultra-high strength steel. The composition of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, includes: C: 0.27%~0.32%, Si: 1.30%~1.60%, Mn: 0.50%~0.80%, Ni: 0.80%~1.20%, Cr: 3.20%~3.60%, Mo: 0.31%~0.60%, W: 0.40%~0.80%, Nb: 0.001%~0.009%, P: ≤0.005%, S: ≤0.001%, with the balance being iron and unavoidable impurities.

[0032] Specifically, in the composition of the aforementioned 1700MPa grade ultra-high toughness and ultra-high strength steel, Cr+Ni≥4.2% and Cr / Ni≥2.8, where Cr and Ni refer to the mass percentage of the corresponding elements.

[0033] Specifically, in the composition of the aforementioned 1700MPa grade ultra-high toughness and ultra-high strength steel, W+Mo≥0.85% and W / Mo≥0.9, where W and Mo refer to the mass percentage of the corresponding elements.

[0034] The following details the function and dosage selection of the components contained in this invention:

[0035] C: Strengthening of lath martensite matrix through interstitial solid solution, with partial desolvation from supersaturated martensite to form Fe during low-temperature tempering. 2.4 Carbon (C) transition carbides form precipitation strengthening compounds and modulate matrix toughness. With increasing carbon content, both tensile strength and yield strength increase, but impact toughness decreases. Excessive C content lowers the Ms point, increases retained austenite and twinned martensite (which reduces toughness), and impairs the weldability of the steel. To achieve a good balance of strength and toughness, the C content should be controlled between 0.27% and 0.32%.

[0036] Si: As one of the main alloying elements in the steel of this invention, the added Si dissolves into the martensitic matrix, improving the strength of the steel through substitution solid solution strengthening. Simultaneously, because Si is insoluble in carbides, it effectively inhibits the formation of cementite, shifting the temper brittleness zone to a higher temperature, thus slightly increasing the tempering temperature of the steel of this invention and enhancing its toughness. However, excessive Si reduces the solubility of elements such as Mo in the steel matrix, leading to residual alloy carbides during quenching and heating, which impairs the toughness of the steel. Therefore, the Si content in this invention is controlled between 1.3% and 1.6%.

[0037] Mn: The addition of Mn can improve the strength of steel while reducing the martensitic transformation temperature and rate, thereby obtaining a small amount of retained austenite and improving toughness. The appropriate addition of Mn helps maintain the stability of the alloy and ensures consistent performance during the preparation process. Taking all factors into consideration, the Mn content in this invention is controlled at 0.5% to 0.8%.

[0038] Cr: As one of the main alloying elements of the steel of this invention, it can effectively improve the hardenability of the steel and enhance its strength through solid solution strengthening. It can also improve the tempering resistance of the steel, so that the steel can obtain excellent comprehensive properties after quenching and tempering. However, too high Cr content reduces the thermal conductivity of the steel, and also lowers the martensitic transformation temperature (Ms) and increases the proportion of twinned martensite. Therefore, the Cr content of this invention is controlled at 3.2% to 3.6%.

[0039] Ni: As one of the main alloying elements in the steel of this invention, Ni effectively increases stacking fault energy and promotes dislocation cross-slip, thereby improving the toughness of the matrix. At the same time, Ni also improves hardenability. However, excessive Ni increases cost, lowers the martensitic transformation temperature (Ms), and increases the proportion of twinned martensite. Therefore, the Ni content in this invention is controlled between 0.8% and 1.2%.

[0040] To achieve strength, toughness, and hardenability comparable to high-alloy ultra-high-strength steel, the total amount of Cr and Ni added in this invention must not be too low. Taking all factors into consideration, the Cr+Ni ratio in this invention needs to be controlled at ≥4.2%. Since Cr has a greater effect on improving hardenability than Ni, and its cost is lower, the Cr / Ni ratio is controlled at ≥2.8.

[0041] Mo and W are both elements in the same group, with similar properties and relatively large sizes. They tend to agglomerate at grain boundaries, preventing impurity aggregation, strengthening grain boundary bonding, and improving toughness. Both form solid solution reinforcement within the matrix and, as carbide-forming elements, contribute to precipitation reinforcement. However, both Mo and W can form M6C carbides with high resolution temperatures, increasing the solution temperature and making dissolution more difficult. High-temperature solution treatment can cause coarse grains, damaging toughness; insufficient solution temperature can leave undissolved large carbide particles, also impairing toughness. Therefore, in this invention, the Mo content is controlled at 0.31%–0.6%, and the W content at 0.4%–0.8%.

[0042] Specifically, to achieve a certain level of strengthening and grain boundary purification effect, this invention controls the sum of W and Mo additions to be ≥0.85%. Furthermore, since W atoms are larger than Mo atoms, their grain boundary segregation ability is stronger, resulting in a more pronounced grain boundary purification effect. Each 0.5% increase in W raises the solution temperature by approximately 50–80°C, while each 0.5% increase in Mo raises it by approximately 100–150°C, meaning that Mo-containing M6C carbides are more difficult to re-dissolve. On the other hand, adding W is more beneficial for improving the strength of steel at high temperatures. Therefore, this invention controls the W / Mo elemental ratio to be ≥0.9.

[0043] Nb: A microalloying element. During quenching heating, a suitable amount of residual nano-sized NbC carbides helps prevent austenite grain growth and refines the martensite size of the laths after quenching. However, excessive Nb content will form large-sized Nb (C / N) inclusions, reducing the toughness of the steel. Therefore, the Nb content should be controlled between 0.001% and 0.009%.

[0044] To further improve the overall performance of the aforementioned 1700MPa grade ultra-high toughness and ultra-high strength steel, the composition of the aforementioned 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, is as follows: C: 0.27%–0.31%, Si: 1.32%–1.55%, Mn: 0.53%–0.75%, Ni: 0.90%–1.20%, Cr: 3.30%–3.58%, Mo: 0.40%–0.60%, W: 0.40%–0.70%, P: ≤0.005%, S: ≤0.001%, Nb: 0.005%–0.009%, with the balance being iron and unavoidable impurities.

[0045] Preferably, in the composition of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, the Cr+Ni content is controlled to be 4.3% to 4.65%, and the Cr / Ni ratio is controlled to be 2.9 to 3.7.

[0046] Preferably, in the composition of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, the W+Mo content is controlled to be 0.85% to 1.15%, and the W / Mo ratio is controlled to be 0.9% to 1.2%.

[0047] It should be noted that in the composition of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, the content of impurity elements is as follows: residual Al content is 0.01% to 0.04%, residual N content is 0.001% to 0.007%, and O content is 0.0007% to 0.0014%.

[0048] Specifically, the yield strength ratio (yield strength / tensile strength) of the aforementioned 1700MPa grade ultra-high toughness ultra-high strength steel is 0.77 to 0.80.

[0049] It should be noted that in the 1700MPa grade ultra-high toughness and ultra-high strength steel of this invention, by precisely controlling the types and contents of alloying elements such as Cr, Ni, Mo, W, and Nb, as well as the synergistic quantitative relationship between different elements, the yield strength ratio of the 1700MPa grade ultra-high toughness and ultra-high strength steel is ensured to be between 0.77 and 0.80. The significance is that, macroscopically, a lower yield strength ratio allows the material to yield earlier and begin the plastic deformation stage, thereby absorbing a large amount of energy and significantly improving toughness indicators such as impact energy and fracture toughness; microscopically, during crack propagation, the area of ​​the plastic zone at the crack tip is significantly larger than that of other steels of the same strength grade. However, on the other hand, excessively low yield strength will also limit the application scenarios of the steel. Therefore, the yield strength ratio of the steel of this invention is controlled within the range of 0.77 to 0.80.

[0050] This invention also provides a method for preparing the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, comprising:

[0051] Step 1: Electric furnace smelting;

[0052] Step 2: Ladle refining and electrode rod casting;

[0053] Step 3: Obtain steel ingots by vacuum consumable remelting or electroslag remelting, and then anneal the steel ingots;

[0054] Step 4: Heating and homogenizing the steel ingot, then forging;

[0055] Step 5: The forging is heat-treated to obtain 1700MPa grade ultra-high toughness and ultra-high strength steel.

[0056] Specifically, in step 1 above, during electric furnace smelting, the oxidation temperature is controlled to be ≥1580℃, the tapping temperature is controlled to be ≥1650℃, and P ≤0.005% before tapping from the electric furnace.

[0057] Specifically, in step 1 above, during the electric furnace smelting process, the final carbon content is controlled to be 0.08% to 0.15%. The oxygen blowing dephosphorization effect is ensured by controlling the final carbon content. When tapping steel from the electric furnace, the Cr and Mn contents are controlled to the lower limit of the specification.

[0058] Specifically, in step 2 above, ladle refining includes: adding 1.0 to 3.0 kg / t of aluminum ingots to the LF ladle according to the mass of the molten steel, adding refining slag and lime to form slag, slag whitening time ≥ 30 min, tapping steel at a temperature ≥ 1620℃, and ensuring S ≤ 0.0012%; total evacuation time of the VD furnace ≥ 35 min, ultimate vacuum holding time ≥ 15 min, and H ≤ 1.5 ppm after evacuation.

[0059] Specifically, in step 2 above, when the LF furnace ladle temperature is ≥1575℃, Si-Ca powder and Al powder are added for reduction, and Ar gas is adjusted to ensure that the molten steel is not exposed during tumbling. After tapping from the LF furnace, 50% of the slag is added with VD, and after sampling and temperature measurement, the furnace is evacuated. The total evacuation time is ≥35min, the ultimate vacuum is maintained for ≥15min, and the H content after evacuation is ≤1.5ppm.

[0060] Specifically, in step 2 above, during the casting of the electrode rod, argon gas is introduced during the casting process, and carbon-free protective slag is added. The casting temperature is controlled at 1520-1560℃. After casting, the mold is cooled for ≥24 hours and then annealed at a temperature of 600-750℃ for ≥25 hours.

[0061] Specifically, in step 2 above, in order to ensure that the remelted steel ingot has a uniform chemical composition and good surface quality, it is preferable to perform riser removal and surface peeling treatment on the steel ingot used as an electrode rod.

[0062] Specifically, in step 3 above, during vacuum consumable remelting or electroslag remelting, the melting rate during the steady-state melting stage is controlled to be 4.5–11.5 kg / min, the voltage to be 21–29 V, and the steady-state melting current to be 6.0–20.0 kA.

[0063] Specifically, in step 3 above, in the vacuum self-consuming remelting or electroslag remelting smelting process, in order to obtain a uniform and dense solidification structure of the steel ingot, the melting rate v (kg / min) in the steady-state melting stage and the ingot diameter (i.e., the diameter of the crystallizer, the diameter of the steel ingot) D (mm) conform to the relationship v = (0.0115~0.0125) × D - 3.

[0064] Specifically, in step 3 above, during vacuum consumable remelting, in order to further improve the cooling effect and avoid the molten pool from deepening significantly as the smelting time increases, helium is started to be charged when the weight of the steel ingot in the crystallizer is ≥10% of the total electrode weight. The helium pressure is 200-700 Pa, for example, 300 Pa, 400 Pa, 500 Pa, or 600 Pa.

[0065] Specifically, in step 3 above, in the electroslag remelting smelting, it is preferred to use a binary slag system of CaF2 and Al2O3 (ratio 70%:30%) for smelting. After the slag is baked, it is hot-sent into the storage tank and protected by argon gas. Argon gas is passed through for 5 to 20 minutes before power is supplied, and argon gas is supplied throughout the entire process after power is supplied.

[0066] Specifically, the above preparation method can be used to prepare large ingots with a diameter of 508 mm or more, such as 660 mm to 1200 mm.

[0067] Specifically, in step 3 above, the steel ingot annealing includes: cooling the steel ingot under vacuum for 1 to 3 hours, breaking the vacuum, and then hot-heat annealing at a temperature of 600 to 750°C for more than 30 hours.

[0068] Specifically, in step 3 above, the annealed steel ingot has a uniform composition and a low degree of microsegregation. For example, the dendrite spacing at half the radius of the feeding end of the annealed steel ingot is 210-350 μm and the Mo element segregation coefficient is 0.69-0.79, while the dendrite spacing in the core is 390-805 μm and the Mo element segregation coefficient is 0.60-0.75.

[0069] It should be noted that, through precise control of the above-mentioned components, combined with precise control of steps and process parameters such as electric furnace melting, ladle refining, electrode rod casting, vacuum consumable remelting or electroslag remelting, and ingot annealing, this invention ensures that the composition of large-sized steel ingots is uniform and the degree of micro-segregation is mild. This avoids the precipitation of large-sized carbides caused by elemental agglomeration, which would damage toughness, and provides a good material basis for forging and heat treatment.

[0070] Specifically, step 4 above includes:

[0071] S41. Place the steel ingot into a heating furnace for heating, and keep it at a temperature of 1150~1280℃;

[0072] S42. Forging: The initial forging temperature of the steel ingot is ≥1100℃, and multiple upsetting and drawing forging are adopted.

[0073] Specifically, in S42 above, after pre-upsetting, three upsettings and three drawings are performed, and finally the billet is formed directly by a fast forging machine or by radial forging by a fast forging machine and a precision forging machine; the amount of reduction in each upsetting is more than 1 / 2 of the original ingot height; the reheating temperature is the same each time, which is 1150~1280℃.

[0074] Specifically, in S42 above, the quick forging integral forging ratio is ≥6, preferably ≥8.

[0075] Specifically, in S42 above, the final forging deformation ratio of the fast forging is ≥2, the final forging temperature range is 780~950℃, and the preferred final forging temperature is 780~900℃.

[0076] Specifically, in S42 above, if rapid forging is followed by radial forging, the initial forging temperature of radial forging must be ≥870℃, and the final forging temperature must be ≥780℃. The final heat of rapid forging must still meet the requirement of a deformation ratio ≥2, and the final forging temperature range is 780~950℃. If reheating is required before radial forging, the heating temperature must not exceed 950℃.

[0077] It should be noted that in step 4 above, the steel of this invention contains carbide-forming elements such as C, W, Mo, and Nb, including M6C containing W and Mo and MC-type carbides containing Nb. These elements need to be fully dissolved to ensure their dispersion and precipitation during subsequent hot working and heat treatment. This prevents the formation of large-sized carbide particles that could damage toughness or prevent alloying elements from entering the matrix and affecting solid solution strengthening. Therefore, the pre-forging heating temperature is set to 1150–1280°C, and the reheating temperature is the same for each reheat.

[0078] It should be noted that, in S42 above, the steel of the present invention is used within a deformation temperature range of 1050℃ to 1250℃ and a deformation time of 0.01 to 0.5s. -1 Within the strain rate range, the energy dissipation rate reaches its peak, and the steel grade exhibits the greatest tendency for dynamic recrystallization. After undergoing severe strain, the deformed grains become polygonized through dislocations, forming dislocation-free low-energy regions as subgrains. These subgrains gradually grow into effective recrystallization nuclei by consuming the surrounding high-dislocation distortion regions, releasing a large amount of deformation-stored energy. Therefore, to achieve sufficient deformation to break the original as-cast structure and achieve adequate homogenization, the initial forging temperature of the steel ingot is set to 1100℃ or higher.

[0079] It should be noted that the steel of this invention is smelted using a low-cost electric furnace + ladle refining + vacuum arc remelting or electroslag remelting process. Compared with the double vacuum process, this reduces the vacuum stage. Al is introduced during the preparation process for deoxidation, resulting in higher residual Al and N content compared to the material produced by the double vacuum process. For example, the residual Al content is 0.01%–0.04%, and the residual N content is 0.001%–0.007%. Under these conditions, AlN typically precipitates above 980°C. If the forging process remains within this temperature range for an extended period without sufficient deformation, the AlN will coarsen into needle-like shapes. Although the final size is only in the hundreds of nanometers, the sharp shape causes severe local stress concentration, significantly impairing toughness. Therefore, to change this situation, on the one hand, reheating in the furnace later allows the formed AlN to be re-dissolved; on the other hand, cooling to a lower temperature below 980°C before deformation will interrupt the formation and coarsening process of AlN, preventing the formation of needle-like AlN. Therefore, since there is no opportunity to remelt the steel during the preparation process of this invention, the final deformation process of this invention must be strictly controlled, and sufficient deformation of ≥2 deformation amount must be carried out at 950°C and below in order to suppress the formation of needle-like AlN.

[0080] It should be noted that, because the deformation in radial forging is relatively small and concentrated in the surface layer, the deformation amount and stopping temperature of the final heat in rapid forging remain crucial for processes following rapid forging. These must be controlled according to the same process requirements as those for processes involving only rapid forging. Furthermore, if the workpiece temperature drops excessively before radial forging, the reheating temperature must be strictly controlled to prevent insufficient deformation leading to excessively coarse grains and AlN precipitation. In this stage, it is safe to control the heating temperature below the upper limit of the quenching temperature and the lower limit of the AlN precipitation temperature.

[0081] Specifically, in step 5 above, the forging undergoes annealing, quenching, and tempering to obtain 1700MPa grade ultra-high strength steel.

[0082] Specifically, in step 5 above, the annealing holding temperature is 600-700℃, the quenching holding temperature is 900-950℃, and the tempering holding temperature is 200-300℃.

[0083] Specifically, in step 5 above, the microstructure obtained after quenching is a lath martensite matrix, a small amount of thin film retained austenite with a width of 10-30 nm and NbC particles with a diameter of 15-30 nm. Tempering precipitates fine and dispersed ε-carbides to obtain a good combination of strength and toughness.

[0084] Specifically, the microstructure of the 1700MPa grade ultra-high toughness and ultra-high strength steel obtained in step 5 above consists of lath martensite matrix, a small amount of thin film retained austenite with a width of 10-30nm and NbC particles with a diameter of 15-30nm, and finely dispersed ε-carbides, wherein the volume fraction of thin film retained austenite is 1%-3%.

[0085] Specifically, the 1700MPa grade ultra-high toughness and ultra-high strength steel of this invention exhibits excellent performance and can meet application requirements. For example, its tensile strength is above 1700MPa (e.g., 1718-1780MPa), its yield strength is above 1350MPa (e.g., 1351-1395MPa), its yield strength ratio is between 0.77 and 0.80, its elongation A is above 12% (e.g., 12.5%-15%), its reduction of area Z is above 52% (e.g., 54%-61%), its impact energy KU2 is above 80J (e.g., 84-110J), and its fracture toughness K... IC 150 MPa·m 1 / 2 Above (e.g., 167–190 MPa·m) 1 / 2 ).

[0086] The advantages of precise control of the composition and process parameters of the steel of the present invention will be demonstrated below with specific embodiments and comparative examples.

[0087] Examples 1-5 of the present invention provide a 1700MPa grade ultra-high toughness and ultra-high strength steel and its preparation method. The chemical composition and impurity element content are shown in Table 1. The preparation method includes the following steps:

[0088] Step 1: Electric furnace smelting;

[0089] Step 2: Ladle refining and electrode rod casting;

[0090] Step 3: Obtain steel ingots by vacuum arc remelting or electroslag remelting, followed by annealing of the steel ingots;

[0091] Step 4: Heating and forging the steel ingot;

[0092] Step 5: The forging is heat-treated to obtain 1700MPa grade ultra-high toughness and ultra-high strength steel.

[0093] Specifically, in step 1, a 20-ton electric furnace is used for smelting. During the electric furnace smelting, the oxidation temperature is controlled at ≥1580℃, the tapping temperature is controlled at ≥1650℃, the P content before tapping is ≤0.005%, and the final carbon content is controlled at 0.08%~0.15%.

[0094] Specifically, in step 2, ladle refining includes: adding 1.0 to 3.0 kg / t of aluminum ingots to the LF ladle according to the mass of the molten steel, adding refining slag and lime to form slag, slag whitening time ≥ 30 min, tapping steel at a temperature ≥ 1620℃, and ensuring S ≤ 0.0012%; total evacuation time of the VD furnace ≥ 35 min, ultimate vacuum holding time ≥ 15 min, and H ≤ 1.5 ppm after evacuation.

[0095] Specifically, in step 2 above, when the LF furnace ladle temperature is ≥1575℃, Si-Ca powder and Al powder are added for reduction, and Ar gas is adjusted to ensure that the molten steel is not exposed during tumbling. After tapping from the LF furnace, 50% of the slag is added with VD, and after sampling and temperature measurement, the furnace is evacuated. The total evacuation time is ≥35min, the ultimate vacuum is maintained for ≥15min, and the H content after evacuation is ≤1.5ppm. During electrode rod casting, argon gas is introduced during tapping and casting, and carbon-free protective slag is added. The casting temperature is controlled at 1520~1560℃. After casting, the mold is cooled for ≥24h and then annealed at 600~750℃ for ≥25h.

[0096] Specifically, in step 3 above, during vacuum consumable remelting or electroslag remelting, the melting rate during the steady-state melting stage is controlled to be 4.5–11.5 kg / min, the voltage to be 21–29 V, and the steady-state melting current to be 6.0–20.0 kA.

[0097] Specifically, in step 3 above, during vacuum arc remelting, helium is introduced into the crystallizer when the weight of the steel ingot is ≥10% of the total electrode weight, and the helium pressure is 200–700 Pa. The steel ingot annealing includes: cooling the steel ingot under vacuum for 1–3 hours, then breaking the vacuum, followed by hot annealing at a temperature of 600–750°C for at least 30 hours.

[0098] Specifically, in step 3 above, the annealed steel ingot has a uniform composition and a low degree of microsegregation. For example, the dendrite spacing at half the radius of the feeding end of the annealed steel ingot is 210-350 μm and the Mo element segregation coefficient is 0.69-0.79, while the dendrite spacing in the core is 390-805 μm and the Mo element segregation coefficient is 0.60-0.75.

[0099] Specifically, step 4 above includes:

[0100] S41. Place the steel ingot into a heating furnace for heating, and keep it at a temperature of 1150~1280℃;

[0101] S42. Forging: The initial forging temperature of the steel ingot is ≥1100℃, and multiple upsetting and drawing forging are adopted.

[0102] Specifically, in S42 above, after pre-upsetting, three upsettings and three drawings are performed, and finally the billet is formed directly by a fast forging machine or by radial forging by a fast forging machine and a precision forging machine; the amount of reduction in each upsetting is more than 1 / 2 of the original ingot height; the reheating temperature is the same each time, which is 1150~1280℃.

[0103] Specifically, in S42 above, the quick forging integral forging ratio is ≥6, preferably ≥8.

[0104] Specifically, in S42 above, the final forging deformation ratio of the fast forging is ≥2, the final forging temperature range is 780~950℃, and the preferred final forging temperature is 780~900℃.

[0105] Specifically, in S42 above, if rapid forging is followed by radial forging, the initial forging temperature of radial forging must be ≥870℃, and the final forging temperature must be ≥780℃. The final heat of rapid forging must still meet the requirement of a deformation ratio ≥2, and the final forging temperature range is 780~950℃. If reheating is required before radial forging, the heating temperature must not exceed 950℃.

[0106] Specifically, in step 5 above, the annealing temperature is 600-700℃, the quenching temperature is 900-950℃, and the tempering temperature is 200-300℃.

[0107] The key process parameters of the embodiments are shown in Table 2.

[0108] The microstructure of the steel in Examples 1-5 consists of a lath martensite matrix, a small amount of thin film retained austenite with a width of 10-30 nm, NbC particles with a diameter of 15-30 nm, and finely dispersed ε-carbides, wherein the volume fraction of thin film retained austenite is 1%-3%.

[0109] The inventors conducted extensive experimental research during the research process, and some unsatisfactory solutions are now presented as comparative examples.

[0110] Comparative Example 1

[0111] This comparative example provides an ultra-high strength steel, the chemical composition of which is shown in Table 1 above, and the preparation method is the same as that in Example 1.

[0112] The main difference between this comparative example and Example 1 lies in the composition design. In this comparative example, Cr+Ni = 3.74% and Cr / Ni = 0.619, which does not meet the requirements of the present invention of Cr+Ni ≥ 4.2% and Cr / Ni ≥ 2.8. Since Ni is more expensive than Cr and has a lower strengthening effect, it is also less effective in improving hardenability than Cr. Therefore, the cost of this comparative example is higher than that of Example 1, and its strength, toughness and hardenability are lower than those of Example 1. The alloy system design has not been optimized, and the performance cannot reach the minimum design value of the present invention.

[0113] Comparative Example 2

[0114] This comparative example provides an ultra-high strength steel, the chemical composition of which is shown in Table 1 above, and the preparation method is the same as that in Example 1.

[0115] The main difference between this comparative example and Example 1 lies in the composition design. This comparative example has a higher Mo content and a lower W content, with W+Mo = 1.0% and W / Mo = 0.129, satisfying the requirement of W+Mo ≥ 0.9% for this invention, but not W / Mo ≥ 0.9. Because the M6C type Mo carbide has a higher dissolution temperature, it cannot be completely dissolved using the same heat treatment regime, resulting in stress concentration of large carbide particles in the matrix, impairing toughness. Simultaneously, the excessive presence of Mo and W in the precipitated phase leads to insufficient solid solution strengthening effect, thereby reducing strength.

[0116] Comparative Example 3

[0117] This comparative example provides an ultra-high strength steel, which comes from the same electric furnace as Example 3. The chemical composition is the same as that of Example 3, and the preparation method is mostly the same as that of Example 3. The difference is that the melting rate of the vacuum self-consuming remelting steady-state melting stage of the same Φ920mm ingot is set at 9.8kg / min, which exceeds the upper limit of the relationship between the melting rate v (kg / min) and the ingot diameter (i.e., the crystallizer diameter, the ingot diameter) D (mm) in the steady-state melting stage specified in this invention, v = (0.0115~0.0125)×D-3, and the helium cooling effect is not correspondingly improved.

[0118] The vacuum consumable remelting process in this comparative example resulted in an excessively deep molten pool, leading to significant microsegregation. The dendrite spacing at half the radius of the feeding end of the consumable ingot was 533 μm, while the dendrite spacing at the core was 820 μm. Mo segregation was severe, and large M6C particles containing W and Mo were visible in the dendrite gaps. Subsequent forging heating and heat treatment processes could not completely eliminate the segregation and fully dissolve the carbides, resulting in a slight decrease in strength and a significant decrease in toughness.

[0119] Comparative Example 4

[0120] This comparative example provides an ultra-high strength steel, derived from the same electric furnace as Example 4, with the same chemical composition and preparation method. Both examples employ an electric furnace + LF + VD + electroslag remelting process, resulting in a higher nitrogen (N) content. The difference between this comparative example and Example 4 is that in Example 4, after forming a Φ450mm bar, it was further reheated to 1180℃ and then forged to Φ340mm using a fast forging process. The final forging deformation ratio was 1.75, and the final forging temperature was 970℃, failing to meet the requirements of a final forging deformation ratio ≥2 and a final forging temperature range of 780–950℃ as specified in this invention. Subsequent direct forging without reheating resulted in a finished product of Φ270mm. This process caused the already relatively high N content in the material to combine with Al, forming AlN, which severely impaired toughness.

[0121] Comparative Example 5

[0122] This comparative example provides an ultra-high strength steel, derived from the same electric furnace as Example 5, with the same chemical composition and preparation method. Both employ an electric furnace + LF + VD + electroslag remelting process, resulting in a higher nitrogen (N) content. In this comparative example, the amount of aluminum ingots added for deoxidation during ladle refining is close to the upper limit, leading to a higher residual Al content in the material. The main difference between this comparative example and Example 5 is that the forged product of this comparative example has a significantly larger cross-section than that of Example 5, a lower overall deformation ratio, and a final forging deformation ratio of 1.3, with a final forging temperature of 980°C. This does not meet the requirements of a final forging deformation ratio ≥2 and a final forging temperature range of 780–950°C as specified in this invention. This process results in coarser grains in the material, where the already relatively high N impurity combines with Al to form AlN, severely impairing toughness.

[0123] Figure 1 The microstructure of Example 1 shows martensitic laths and thin austenite film.

[0124] Figure 2 The microstructure of Example 1 shows the main reinforcing phase ε-carbide; Figure 3 The strain field at the crack tip (plastic deformation zone) during the fracture toughness test in Example 2 is shown. Figure 4 The microstructure of Example 5 (nanoscale TiC can be seen); Figure 5 As shown in the microstructure of Comparative Example 3, large M6C particles can be observed. Figure 6 For comparative example 4, the stress field at the crack tip (plastic deformation zone) during fracture toughness testing is shown. Figure 7 The microstructure of Comparative Example 5 (containing needle-like AlN) is shown.

[0125] Table 1 Chemical composition, wt%

[0126]

[0127] Note: P: ≤0.005%, S: ≤0.001%.

[0128] Table 2 Key Process Parameters

[0129]

[0130]

[0131]

[0132]

[0133] The main performance test results of the steels in Examples 1-5 and Comparative Examples 1-5 are shown in Table 3.

[0134] Table 3 Mechanical Properties

[0135]

[0136] The above description is only a preferred embodiment of the present invention, but the scope of protection of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by those skilled in the art within the scope of the technology disclosed in the present invention should be included within the scope of protection of the present invention.

Claims

1. A 1700MPa grade ultra-high toughness and ultra-high strength steel, characterized in that, The composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, includes: C: 0.27%~0.32%, Si: 1.30%~1.60%, Mn: 0.50%~0.80%, Ni: 0.80%~1.20%, Cr: 3.20%~3.60%, Mo: 0.31%~0.60%, W: 0.40%~0.80%, Nb: 0.001%~0.009%, P: ≤0.005%, S: ≤0.001%, with the balance being iron and unavoidable impurities; The W+Mo ratio is 0.85%~1.15%, and the W / Mo ratio is 0.9~1.2; where W and Mo refer to the mass percentage of the corresponding elements. The preparation method of the 1700MPa grade ultra-high toughness and ultra-high strength steel includes: electric furnace melting; ladle refining; electrode rod casting; vacuum consumable remelting or electroslag remelting to obtain steel ingots; steel ingot annealing; steel ingot heating homogenization treatment; forging; heat treatment; In the vacuum self-consuming remelting or electroslag remelting, the melting rate v in the steady-state melting stage conforms to the following relationship with the ingot diameter D: v = (0.0115~0.0125) × D - 3; Forging includes rapid forging, where the final forging deformation ratio is ≥2 and the final forging temperature range is 780~950℃; The microstructure of the 1700MPa grade ultra-high toughness and ultra-high strength steel consists of a lath martensite matrix, a small amount of thin film retained austenite with a width of 10~30nm, NbC particles with a diameter of 15~30nm, and finely dispersed ε-carbides, wherein the volume fraction of thin film retained austenite is 1%~3%; The 1700MPa grade ultra-high toughness ultra-high strength steel has a tensile strength of over 1700MPa, a yield strength of over 1350MPa, an impact energy (KU2) of 84~110J, and a fracture toughness (KIC) of 167~190MPa·m. 1 / 2 .

2. The 1700MPa grade ultra-high toughness and ultra-high strength steel according to claim 1, characterized in that, The composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel is Cr+Ni≥4.2% and Cr / Ni≥2.8, where Cr and Ni refer to the mass percentage of the corresponding elements.

3. The 1700MPa grade ultra-high toughness and ultra-high strength steel according to claim 1, characterized in that, The composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel has a W / Mo ratio of 0.9~1.

1.

4. The 1700MPa grade ultra-high toughness and ultra-high strength steel according to claim 1, characterized in that, The composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel, by mass percentage, includes: C: 0.27%~0.31%, Si: 1.32%~1.55%, Mn: 0.53%~0.75%, Ni: 0.90%~1.20%, Cr: 3.30%~3.58%, Mo: 0.40%~0.60%, W: 0.40%~0.70%, P: ≤0.005%, S: ≤0.001%, Nb: 0.005%~0.009%, with the balance being iron and unavoidable impurities.

5. The 1700MPa grade ultra-high toughness ultra-high strength steel according to any one of claims 1 to 4, characterized in that, The yield strength ratio of the 1700MPa grade ultra-high toughness ultra-high strength steel is 0.77~0.

80.

6. A method for preparing 1700MPa grade ultra-high toughness and ultra-high strength steel according to any one of claims 1 to 5, characterized in that, include: Electric furnace smelting; Ladle refining, electrode rod casting; vacuum consumable remelting or electroslag remelting to obtain steel ingots, steel ingot annealing; steel ingot heating homogenization treatment, forging; Heat treatment; In the vacuum self-consuming remelting or electroslag remelting, the melting rate v in the steady-state melting stage conforms to the following relationship with the ingot diameter D: v = (0.0115~0.0125) × D - 3; Forging includes rapid forging, with a final forging deformation ratio ≥2 and a final forging temperature range of 780~950℃.

7. The preparation method according to claim 6, characterized in that, The final forging temperature range is 780~900℃.

8. The preparation method according to any one of claims 6 to 7, characterized in that, Heat treatment includes annealing, quenching, and tempering. The annealing holding temperature is 600~700℃, the quenching holding temperature is 900~950℃, and the tempering holding temperature is 200~300℃.

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