1700MPa-grade ultrahigh-toughness ultrahigh-strength steel and preparation method thereof

Abstract

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

Description

Technical Field

[0001] The present invention relates to the technical field of ultra-high strength steel, and in particular to a 1700MPa grade ultra-high toughness ultra-high strength steel and a preparation method thereof. Background Art

[0002] Ultra-high-strength steel has been widely used in key load-bearing components such as aircraft landing gear, engine shafts, gears, etc. for many years due to its excellent specific strength and excellent strength and toughness combination. In recent years, with the development of high-performance and low-cost equipment, the market is very sensitive to the cost of raw materials. Traditional secondary hardening steel and maraging steel that rely on high Co / Ni alloy content and ultra-pure smelting process to achieve ultra-high strength and high toughness have been limited in promotion and application. In this context, it is of great significance to achieve a strength and toughness level basically equivalent to that of high-alloy ultra-high-strength steel by reasonably designing smelting, hot working and other process parameters under the composition system of medium and low alloy ultra-high-strength steel. Summary of the invention

[0003] In view of the above situation, the present invention aims to provide a 1700MPa grade ultra-high toughness ultra-high strength steel and a preparation method thereof, which is used to solve at least one of the following problems: the existing ultra-high strength steel has a high alloy element content, high cost, and poor strength-toughness matching.

[0004] The purpose of the present invention is mainly achieved through the following technical solutions:

[0005] On the one hand, the present invention provides a 1700MPa grade ultra-high toughness and ultra-high strength steel. The components of the 1700MPa grade ultra-high toughness and ultra-high strength steel include, by mass percentage: 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%, and the balance is iron and unavoidable impurities.

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

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

[0008] Furthermore, the components of 1700MPa grade ultra-high toughness and ultra-high strength steel include, by mass percentage: 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%, and the balance is iron and unavoidable impurities.

[0009] Furthermore, the microstructure of the 1700MPa grade ultra-high toughness and ultra-high strength steel is a lath martensite matrix + a small amount of film retained austenite and NbC particles, and fine 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 smelting; refining outside the furnace, electrode rod casting; vacuum consumable remelting smelting or electroslag remelting smelting to obtain steel ingots, annealing the steel ingots; heating and homogenizing the steel ingots, forging; and heat treatment.

[0012] Furthermore, in vacuum consumable remelting or electroslag remelting, the melting rate v in the steady-state smelting stage and the ingot diameter D conform to the following relationship: v = (0.0115-0.0125) × D-3.

[0013] Furthermore, the forging includes rapid forging, the final fire deformation ratio of the rapid forging is ≥2, and the final forging temperature ranges from 780 to 950°C.

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

[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 the alloy elements 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, thereby achieving the effect of preparing ultra-high strength steel at low raw material cost.

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

[0018] c) The preparation method of the present invention achieves excellent performance of the steel by precisely controlling process steps and parameters such as homogenization treatment and forging, avoiding the precipitation of needle-shaped harmful phases during forging that damages toughness, thereby ensuring the combination of ultra-high strength and ultra-high toughness.

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

[0020] Other features and advantages of the present invention will be described in the following description, and part of them will become obvious from the description, or will be understood by practicing the present invention. The purpose and other advantages of the present invention can be realized and obtained by the contents particularly pointed out in the written description and the accompanying drawings. BRIEF DESCRIPTION OF THE DRAWINGS

[0021] The drawings are only for the purpose of illustrating particular embodiments and are not to be considered limiting of the present invention. Like reference symbols denote like components throughout the drawings.

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

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

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

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

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

[0027] Figure 6 The stress field at the crack tip (plastic deformation zone) during loading of the fracture toughness test of comparative example 4;

[0028] Figure 7 This is the microstructure of Comparative Example 5 (needle-shaped AlN exists). DETAILED DESCRIPTION

[0029] Preferred embodiments of the present invention will be described in detail below in conjunction with the accompanying drawings, wherein the accompanying drawings constitute a part of the present invention and are used to explain the principles of the present invention together with the embodiments of the present invention.

[0030] Ultra-high strength steel is very sensitive to impurities or second phases in the material and the microcracks caused by them. In order to improve the cleanliness of the material, avoid large-sized precipitation phases and thus improve toughness, high-alloy ultra-high strength steel is mostly smelted using the expensive "double vacuum" (vacuum induction + vacuum consumable remelting) process. Generally, in order to control the overall cost, medium and low alloy ultra-high strength steel is smelted using a lower-cost process of electric furnace + refining outside the furnace + vacuum consumable / electroslag remelting. This requires, on the one hand, fine control of the refining process to obtain high-cleanliness steel ingots, and on the other hand, optimized design of the hot processing process to obtain high-quality original forging structure.

[0031] The present invention provides a 1700MPa grade ultra-high toughness and ultra-high strength steel. The components of the 1700MPa grade ultra-high toughness and ultra-high strength steel include, by mass percentage, 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%, and the balance is iron and unavoidable impurities.

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

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

[0034] The following is a detailed description of the effects and dosage of the components in the present invention:

[0035] C: The strength of the lath martensite matrix is ​​improved by interstitial solid solution strengthening, and Fe is partially dissolved from the supersaturated martensite during low-temperature tempering. 2.4 C transition carbide forms precipitation strengthening and adjusts matrix toughness. As the carbon content increases, both tensile strength and yield strength increase, but impact toughness decreases. Too high a C content will reduce the Ms point, increase the residual austenite and the twin martensite that reduces toughness, and damage the welding performance of the steel. In order to achieve a good combination of strength and toughness, the C content should be controlled at 0.27% to 0.32% after comprehensive consideration.

[0036] Si: As one of the main alloying elements of the steel of the present invention, the added Si dissolves into the martensite matrix, and improves the strength of the steel through substitution solid solution strengthening. At the same time, because Si is insoluble in carbides, it effectively inhibits the formation of cementite, pushes the tempering brittleness zone to a higher temperature, and makes the tempering temperature of the steel of the present invention slightly increased, thereby enhancing the toughness. However, too high Si reduces the solubility of elements such as Mo in the steel matrix, resulting in residual alloy carbides during quenching and heating, which damages the toughness of the steel. Therefore, the Si content of the present invention is controlled between 1.3% and 1.6%.

[0037] Mn: Mn addition can improve the strength of steel, while reducing the martensitic phase transition temperature and phase transition speed to obtain a small amount of residual austenite and improve toughness. The addition of an appropriate amount of Mn helps maintain the stability of the alloy and ensures consistent performance during the preparation process. Taking all factors into consideration, in the present invention, the Mn content is controlled at 0.5% to 0.8%.

[0038] Cr: As one of the main alloying elements of the steel of the present invention, Cr can effectively improve the hardenability of the steel, and improve the strength of the steel by solid solution strengthening, and also improve the tempering resistance of the steel, so that the steel obtains excellent comprehensive properties after quenching and tempering; but too high Cr content reduces the thermal conductivity of the steel, and also has the problem of reducing the martensitic phase transformation temperature (Ms) and increasing the proportion of twin martensite, so the Cr content of the present invention is controlled at 3.2% to 3.6%.

[0039] Ni: As one of the main alloying elements of the steel of the present invention, Ni effectively increases the stacking fault energy, promotes dislocation cross slip, thereby improving the matrix toughness, and at the same time Ni has the effect of improving hardenability; but too high Ni increases the cost, reduces the martensitic phase transition temperature (Ms), and increases the proportion of twin martensite. Therefore, the Ni content of the present invention is controlled between 0.8% and 1.2%.

[0040] In order to achieve toughness and hardenability close to that of high-alloy ultra-high strength steel, the total amount of Cr and Ni elements added in the present invention should not be too low. Taking all factors into consideration, the Cr+Ni ratio should be controlled to be ≥ 4.2%. However, the effect of Cr on improving hardenability is greater than that of Ni, and its cost is lower, so the Cr / Ni ratio is controlled to be ≥ 2.8.

[0041] Mo and W: They are elements of the same family, with similar properties and larger sizes. They are easily segregated on the grain boundaries to avoid impurity segregation, enhance grain boundary bonding, and improve toughness. They are solid-dissolved in the matrix to form solid solution strengthening. As carbide-forming elements, they can also contribute to precipitation strengthening. However, both Mo and W can form M6C carbides with high resolubility temperatures, which increases the solid solution temperature and the difficulty of dissolution. If high-temperature solid solution is used, coarse grains will damage toughness. If the solid solution temperature is not enough, residual undissolved large particles of carbides will also damage toughness. Therefore, the Mo content of the present invention is controlled at 0.31% to 0.6%, and the W content is controlled at 0.4% to 0.8%.

[0042] In particular, in order to form a certain amount of strengthening effect and grain boundary purification effect, the present invention controls the sum of the addition amounts of W and Mo to be ≥0.85%. In addition, since W atoms are larger in volume than Mo atoms, their grain boundary segregation ability is stronger, the grain boundary purification effect is stronger, and each increase of 0.5% W increases the solid solution temperature by about 50 to 80°C, while each increase of 0.5% Mo increases the solid solution temperature by about 100 to 150°C, that is, the M6C carbide containing Mo is more difficult to dissolve back; on the other hand, adding W is more beneficial to improving the strength of steel at high temperatures. Therefore, the present invention controls the W / Mo element ratio to be ≥0.9.

[0043] Nb: Microalloying element. During quenching and heating, a proper amount of residual nano-scale NbC carbides are formed to prevent the growth of austenite grains and refine the size of lath martensite after quenching. However, too high Nb content will form large-sized Nb (C / N) inclusions, reducing the toughness of the steel. Therefore, the Nb content is controlled to be 0.001% to 0.009%.

[0044] In order to further improve the comprehensive performance of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel, the components of the above-mentioned 1700MPa grade ultra-high toughness and ultra-high strength steel are calculated by mass percentage: 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%, and the balance is iron and unavoidable impurities.

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

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

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

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

[0049] It should be noted that in the 1700MPa-level ultra-high toughness and ultra-high strength steel of the present invention, the types and contents of the alloying elements Cr, Ni, Mo, W, and Nb, as well as the synergistic quantitative relationship of different elements, ensure that the yield ratio of the 1700MPa-level ultra-high toughness and ultra-high strength steel is between 0.77 and 0.80. Its significance lies in that, from a macroscopic point of view, a lower yield ratio can make the material yield earlier and start the plastic deformation stage, thereby absorbing a large amount of energy, greatly improving toughness indicators such as impact energy and fracture toughness; from a microscopic point of view, during the crack propagation process, the plastic zone area at the crack tip is significantly larger than other steels of the same strength level. On the other hand, too low a yield strength will also limit the application scenarios of steel. Therefore, the yield ratio of the steel of the present invention is controlled to be in the range of 0.77 to 0.80.

[0050] The present 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: refining outside the furnace and pouring electrode rods;

[0053] Step 3, vacuum consumable remelting or electroslag remelting to obtain a steel ingot, and annealing the steel ingot;

[0054] Step 4: heating and homogenizing the steel ingot and forging it;

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

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

[0057] Specifically, in the above step 1, during the electric furnace smelting process, the endpoint carbon content is controlled to be 0.08% to 0.15%, and the oxygen blowing and phosphorus removal effect is ensured by the endpoint carbon control. When the electric furnace is tapped, the Cr and Mn contents are controlled to be the lower limits of the specifications.

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

[0059] Specifically, in the above step 2, when the ladle temperature of the LF furnace is ≥1575°C, Si-Ca powder and Al powder are added for reduction, and the Ar gas is adjusted to ensure that the molten steel is not exposed when it is turned over. After LF steel is tapped, 50% of the slag is turned over and transferred into VD, and the vacuum is evacuated after sampling and temperature measurement. The total evacuation time is ≥35min, the ultimate vacuum holding time is ≥15min, and the H content after breaking the vacuum is ≤1.5ppm.

[0060] Specifically, in the above step 2, during the pouring of the electrode rod, argon is filled during the steel pouring, carbon-free protective slag is added, the pouring temperature is controlled at 1520-1560°C, the mold is cooled for ≥24h after pouring, and then annealed at 600-750°C for ≥25h.

[0061] Specifically, in the above step 2, in order to ensure that the chemical composition of the remelted steel ingot is uniform and the surface quality is good, the steel ingot used as the electrode rod is preferably subjected to riser removal treatment and surface peeling treatment.

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

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

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

[0065] Specifically, in the above step 3, in the electroslag remelting smelting, it is preferred to use CaF2 and Al2O3 binary slag system (ratio 70%:30%) for smelting. After the slag is baked, it is hot sent to the storage tank and filled with argon for protection; argon is passed for 5 to 20 minutes before power is supplied, and argon is filled throughout the process after power is supplied.

[0066] Specifically, the above preparation method can be used to prepare large ingots with a size of Φ508 mm or larger, such as Φ660 mm to Φ1200 mm.

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

[0068] Specifically, in the above step 3, the composition of the annealed steel ingot is uniform and the degree of micro-segregation is light. For example, the dendrite spacing at the 1 / 2 radius of the feeding end of the annealed steel ingot is 210-350 μm, the Mo element segregation coefficient is 0.69-0.79, the dendrite spacing at the core is 390-805 μm, and the Mo element segregation coefficient is 0.60-0.75.

[0069] It should be noted that the present invention ensures uniform composition of large-sized steel ingots and light micro-segregation through precise control of the above-mentioned components, combined with precise control of electric furnace smelting, refining outside the furnace, electrode rod casting, vacuum consumable remelting or electroslag remelting, ingot annealing and other steps and process parameters, thereby avoiding the precipitation of large-sized carbides that damage toughness due to element segregation, and providing a good material basis for forging and heat treatment.

[0070] Specifically, the above step 4 includes:

[0071] S41, placing the steel ingot into a heating furnace for heating at a holding temperature of 1150-1280°C;

[0072] S42, forging, the steel ingot forging temperature ≥ 1100 ℃, using multiple upsetting and drawing forging.

[0073] Specifically, in the above S42, three-upsetting and three-drawing are performed after pre-upsetting, and finally the forging blank is directly formed by a fast forging machine or radial forging by a fast forging machine + a fine forging machine to obtain the forging blank; the amount of downsetting each time is more than 1 / 2 of the original ingot height; the re-firing temperature each time is the same, which is 1150-1280°C.

[0074] Specifically, in the above S42, the overall forging ratio of the rapid forging is ≥6, preferably ≥8.

[0075] Specifically, in the above S42, the deformation ratio of the last rapid forging is ≥2, and the final forging temperature ranges from 780 to 950°C, preferably from 780 to 900°C.

[0076] Specifically, in the above S42, if the radial forging is switched after the rapid forging, the initial forging temperature of the radial forging is ≥870°C, and the final forging temperature is ≥780°C. The last fire of the rapid forging still needs to meet the deformation ratio ≥2, and the final forging temperature ranges from 780 to 950°C. If the furnace needs to be reheated before the radial forging, the heating temperature shall not exceed 950°C.

[0077] It should be noted that in the above step 4, the steel of the present invention contains carbide-forming elements such as C, W, Mo, and Nb, including M6C containing W and Mo and MC-type carbides containing Nb, etc., which need to be fully dissolved back to disperse and precipitate during subsequent hot working and heat treatment to avoid the formation of large-sized carbide particles that damage toughness or the inability of alloy elements to enter the matrix to affect the solid solution strengthening effect. Therefore, the heating temperature before forging is set to 1150-1280°C, and the re-firing temperature is the same each time.

[0078] It should be noted that in the above S42, the steel of the present invention is within the deformation temperature range of 1050°C to 1250°C and 0.01 to 0.5s -1 Within the strain rate range, the energy dissipation rate reaches the peak value, and the dynamic recrystallization tendency of this steel grade is the largest. After the deformed grains experience severe strain, they are polygonized by dislocations to form low-energy areas without dislocations as subgrains. Subgrains gradually grow into effective recrystallization cores by consuming the surrounding high-dislocation distortion areas, releasing a large amount of deformation storage energy. Therefore, in order to fully deform and break the original cast structure and fully homogenize it, the ingot forging temperature is set to 1100℃ or above.

[0079] It should be noted that the steel of the present invention is smelted by a low-cost electric furnace + refining outside the furnace + vacuum consumable remelting or electroslag remelting process. Compared with the double vacuum process, the vacuum link is reduced. Al needs to be introduced for deoxidation during the preparation process, resulting in the final residual Al and N content being higher than that of the double vacuum process material, such as the residual Al content of 0.01% to 0.04%, and the residual N content of 0.001% to 0.007%. Under this condition, AlN can usually precipitate above 980°C. If it stays in this temperature range for a long time during the forging process and is not sufficiently deformed, AlN will coarsen and become needle-shaped. Although the final size is only a hundred nanometers, the sharp shape causes severe local stress concentration, which will seriously damage the toughness; therefore, to change this situation, on the one hand, if it is reheated in the later stage, the formed AlN can be dissolved again; on the other hand, if it is cooled to a lower temperature below 980°C and then deformed, the formation and coarsening process of AlN will be interrupted, and needle-shaped AlN will not appear. Therefore, since there is no chance to return to the furnace during the preparation process of the present invention, the steel of the present invention needs to strictly control the last fire deformation process and perform sufficient deformation of ≥2 deformation amounts at 950°C or below to inhibit the formation of needle-shaped AlN.

[0080] It should be noted that, since the deformation of radial forging is small and concentrated on the surface, for the process of rapid forging followed by radial forging, the deformation of the last rapid forging and the stopping deformation temperature are still very important and need to be controlled in accordance with the same process requirements as the process flow of only rapid forging. At the same time, if the temperature drop of the workpiece before radial forging is too large, the supplementary heating temperature must be strictly controlled to avoid insufficient deformation causing excessive coarsening of grains and precipitation of AlN. In this link, it is safe to control the heating temperature not to exceed the upper limit of the quenching temperature and the lower limit of the AlN precipitation temperature.

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

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

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

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

[0085] Specifically, the 1700MPa grade ultra-high toughness ultra-high strength steel of the present invention has excellent performance and can meet the use requirements. For example, the tensile strength is above 1700MPa (for example, 1718-1780MPa), the yield strength is above 1350MPa (for example, 1351-1395MPa), the yield strength ratio is between 0.77 and 0.80, the elongation A is above 12% (for example, 12.5%-15%), the cross-sectional shrinkage Z is above 52% (for example, 54%-61%), the impact energy KU2 is above 80J (for example, 84-110J), and the fracture toughness K IC 150MPa·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 are demonstrated below with specific examples and comparative examples.

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

[0088] Step 1, electric furnace smelting;

[0089] Step 2: refining outside the furnace and pouring electrode rods;

[0090] Step 3, vacuum consumable remelting or electroslag remelting to obtain a steel ingot, and annealing the steel ingot;

[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 to be ≥1580°C, the steel-out temperature is ≥1650°C, P is ≤0.005% before steel-out from the electric furnace, and the terminal carbon content is controlled to be 0.08% to 0.15%.

[0094] Specifically, in step 2, the refining outside the furnace includes: adding 1.0-3.0 kg / t of aluminum ingots according to the mass of molten steel into the LF ladle, adding refining slag and lime to make slag, the slag whitening time is ≥30 min, the temperature is ≥1620°C for steel tapping, and S≤0.0012% is ensured; the total evacuation time of the VD furnace is ≥35 min, the ultimate vacuum holding time is ≥15 min, and H≤1.5 ppm after breaking the vacuum.

[0095] Specifically, in the above step 2, when the temperature of the LF furnace ladle is ≥1575°C, Si-Ca powder and Al powder are added for reduction, and the Ar gas is adjusted to ensure that the molten steel is not exposed when it is turned over. After LF steel is tapped, 50% of the slag is transferred into VD, and the vacuum is evacuated after sampling and temperature measurement. The total vacuum time is ≥35min, the ultimate vacuum holding time is ≥15min, and the H content after breaking the vacuum is ≤1.5ppm. During the pouring of the electrode rod, argon is filled during steel pouring, and carbon-free protective slag is added. The pouring temperature is controlled at 1520-1560°C. After pouring, the mold is cooled for ≥24h and then annealed. The annealing temperature is 600-750°C and the insulation time is ≥25h.

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

[0097] Specifically, in the above step 3, during the vacuum consumable remelting, when the weight of the steel ingot in the crystallizer is ≥10% of the total electrode weight, helium is filled, and the helium pressure is 200-700 Pa. The steel ingot annealing includes: the steel ingot is cooled under vacuum for 1-3 hours, then broken, hot annealing, the annealing temperature is 600-750°C, and the holding time is more than 30 hours.

[0098] Specifically, in the above step 3, the composition of the annealed steel ingot is uniform and the degree of micro-segregation is light. For example, the dendrite spacing at the 1 / 2 radius of the feeding end of the annealed steel ingot is 210-350 μm, the Mo element segregation coefficient is 0.69-0.79, the dendrite spacing at the core is 390-805 μm, and the Mo element segregation coefficient is 0.60-0.75.

[0099] Specifically, the above step 4 includes:

[0100] S41, placing the steel ingot into a heating furnace for heating at a holding temperature of 1150-1280°C;

[0101] S42, forging, the steel ingot forging temperature ≥ 1100 ℃, using multiple upsetting and drawing forging.

[0102] Specifically, in the above S42, three-upsetting and three-drawing are performed after pre-upsetting, and finally the forging blank is directly formed by a fast forging machine or radial forging by a fast forging machine + a fine forging machine to obtain the forging blank; the amount of downsetting each time is more than 1 / 2 of the original ingot height; the re-firing temperature each time is the same, which is 1150-1280°C.

[0103] Specifically, in the above S42, the overall forging ratio of the rapid forging is ≥6, preferably ≥8.

[0104] Specifically, in the above S42, the deformation ratio of the last rapid forging is ≥2, and the final forging temperature ranges from 780 to 950°C, preferably from 780 to 900°C.

[0105] Specifically, in the above S42, if the radial forging is switched after the rapid forging, the initial forging temperature of the radial forging is ≥870°C, and the final forging temperature is ≥780°C. The last fire of the rapid forging still needs to meet the deformation ratio ≥2, and the final forging temperature ranges from 780 to 950°C. If the furnace needs to be reheated before the radial forging, the heating temperature shall not exceed 950°C.

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

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

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

[0109] The inventor has conducted a large number of experimental studies during the research process, and now uses some bad solutions 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 of 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%, Cr / Ni=0.619, which do not meet the requirements of the present invention of Cr+Ni≥4.2%, Cr / Ni≥2.8. Since the cost of Ni element is more expensive than that of Cr element, and the strengthening effect is lower than that of Cr element, and the effect of improving hardenability is not as good as that of Cr element, the cost of this comparative example is higher than that of Example 1, and the toughness and hardenability are lower than those of Example 1. The alloy system design is not 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 of Example 1.

[0115] The main difference between this comparative example and Example 1 is the composition design, wherein the Mo content of this comparative example is higher and the W content is lower, W+Mo=1.0%, W / Mo=0.129, which meets the requirement of W+Mo≥0.9% of the present invention, but does not meet W / Mo≥0.9. Since the M6C carbide of Mo has a higher dissolution temperature, it cannot be completely dissolved using the same heat treatment system, causing large-particle carbide to generate stress concentration in the matrix and damage toughness. At the same time, too much Mo and W exist in the precipitation phase, resulting in insufficient solid solution strengthening effect, resulting in reduced strength.

[0116] Comparative Example 3

[0117] This comparative example provides an ultra-high strength steel, which comes from the same electric furnace mother furnace as Example 3, has the same chemical composition as Example 3, and is prepared in a mostly the same manner as Example 3, with the difference that the melting rate in the steady-state melting stage of the vacuum consumable remelting of the same Φ920 mm ingot is set to 9.8 kg / min, which exceeds the upper limit of the relationship between the melting rate v (kg / min) in the steady-state melting stage and the ingot diameter (i.e., crystallizer diameter, ingot diameter) D (mm) specified in the present invention, v = (0.0115 to 0.0125) × D-3, and the helium cooling effect is not improved accordingly.

[0118] In this comparative example, the molten pool in the vacuum consumable remelting process is too deep, resulting in heavier micro-segregation. The dendrite spacing at the 1 / 2 radius of the consumable ingot feeding end is 533μm, and the dendrite spacing in the center is 820μm. The Mo element is severely segregated, and large particles of M6C containing W and Mo can be seen precipitating in the dendrite gaps. The subsequent forging heating and heat treatment processes cannot 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, which comes from the same electric furnace mother furnace as Example 4, has the same chemical composition as Example 4, and the preparation method is mostly the same as Example 4. Since the smelting process of "electric furnace + LF + VD + electroslag remelting" is adopted, the impurity element N content is relatively high. The difference between this comparative example and Example 4 is that after the Φ450mm bar has been formed in Example 4, it is further returned to the furnace for heating at 1180℃, and then it is taken out of the furnace and forged into Φ340mm on the fast forging. The last fire deformation ratio is 1.75, and the final forging temperature is 970℃, which does not meet the requirements of the present invention that the last fire deformation ratio of the fast forging is ≥2, and the final forging temperature range is 780~950℃. There is no subsequent heat supplementation and direct forging, and the finished product specification is Φ270mm. This process causes the relatively high impurity N in the material to combine with Al to form AlN, which seriously damages the toughness.

[0121] Comparative Example 5

[0122] This comparative example provides an ultra-high strength steel, which comes from the same electric furnace mother furnace as Example 5, has the same chemical composition as Example 5, and the preparation method is mostly the same as Example 5. Since the smelting process of "electric furnace + LF + VD + electroslag remelting" is adopted, the impurity element N content is relatively high. In the refining process outside the furnace, the amount of aluminum ingots added for deoxidation in this comparative example is close to the upper limit, resulting in a high content of residual Al in the material. The main difference between this comparative example and Example 5 is that the cross-section of the forged product of this comparative example is significantly larger than that of Example 5, the overall deformation ratio is lower than that of Example 5, and the last fire deformation ratio is 1.3, and the final forging temperature is 980°C, which does not meet the requirements of the present invention that the last fire deformation ratio of fast forging is ≥2, and the final forging temperature range is 780-950°C. This process results in coarse grains of the material, in which the relatively high impurity N combines with Al to form AlN, which seriously damages the toughness.

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

[0124] Figure 2 The microstructure of Example 1 shows the main strengthening phase ε-carbide; Figure 3 The strain field (plastic deformation zone) at the crack tip during loading of the fracture toughness test in Example 2; Figure 4 The microstructure of Example 5 (nanoscale TiC can be seen); Figure 5 This is the microstructure of Comparative Example 3, where it can be seen that large particles of M6C exist; Figure 6 The stress field at the crack tip (plastic deformation zone) during loading of the fracture toughness test of comparative example 4; Figure 7 This is the microstructure of Comparative Example 5 (needle-shaped AlN exists).

[0125] Table 1 Chemical composition, wt%

[0126]

[0127] Remarks: P:≤0.005%, S:≤0.001%.

[0128] Table 2 Key process parameters

[0129]

[0130]

[0131]

[0132]

[0133] The main performance test results of the steel of 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 specific implementation manner of the present invention, but the protection scope of the present invention is not limited thereto. Any changes or substitutions that can be easily conceived by any technician familiar with the technical field within the technical scope disclosed by the present invention should be covered within the protection scope of the present invention.

Claims

1. A 1700MPa grade ultra-high toughness and ultra-high strength steel, characterized in that: The components of the 1700MPa grade ultra-high toughness and ultra-high strength steel include, by mass percentage, 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 remainder is iron and unavoidable impurities.

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

3. The 1700MPa grade ultra-high toughness and ultra-high strength steel according to claim 1, characterized in that: In the composition of the 1700MPa grade ultra-high toughness and ultra-high strength steel, W+Mo≥0.85%, W / Mo≥0.9, wherein W and Mo refer to the mass percentages of the corresponding elements.

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

5. The 1700MPa grade ultra-high toughness and ultra-high strength steel according to any one of claims 1 to 4, characterized in that: The microstructure of the 1700MPa grade ultra-high toughness and ultra-high strength steel is a lath martensite matrix + a small amount of film residual austenite and NbC particles, and fine dispersed ε-carbides.

6. The 1700MPa grade ultra-high toughness and 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 and ultra-high strength steel is 0.77 to 0.

80.

7. A method for preparing the 1700MPa grade ultra-high toughness and ultra-high strength steel according to any one of claims 1 to 6, characterized in that: include: Electric furnace smelting; Refining outside the furnace, pouring of electrode rods; vacuum consumable remelting or electroslag remelting to obtain steel ingots, annealing of steel ingots; heating and homogenization of steel ingots, and forging; Heat treatment.

8. The preparation method according to claim 7, characterized in that: In the vacuum consumable remelting smelting or electroslag remelting smelting, the melting rate v in the steady-state melting stage and the ingot diameter D meet the following relationship: v=(0.0115-0.0125)×D-3.

9. The preparation method according to claim 7, characterized in that: Forging includes rapid forging, the final fire deformation ratio of rapid forging is ≥2, and the final forging temperature ranges from 780 to 950°C.

10. The preparation method according to any one of claims 7 to 9, characterized in that: Heat treatment includes annealing, quenching and tempering; the annealing holding temperature is 600-700°C, the quenching holding temperature is 900-950°C, and the tempering holding temperature is 200-300°C.

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