WO2019222988A1 - 一种屈服强度1100MPa级超细晶高强钢板及其制造方法 - Google Patents

一种屈服强度1100MPa级超细晶高强钢板及其制造方法 Download PDF

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WO2019222988A1
WO2019222988A1 PCT/CN2018/088336 CN2018088336W WO2019222988A1 WO 2019222988 A1 WO2019222988 A1 WO 2019222988A1 CN 2018088336 W CN2018088336 W CN 2018088336W WO 2019222988 A1 WO2019222988 A1 WO 2019222988A1
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temperature
strength
rolling
steel sheet
steel plate
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PCT/CN2018/088336
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English (en)
French (fr)
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闫强军
姜在伟
温长飞
邓想涛
王昭东
吴俊平
张仪杰
王思聪
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南京钢铁股份有限公司
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Priority to PCT/CN2018/088336 priority Critical patent/WO2019222988A1/zh
Publication of WO2019222988A1 publication Critical patent/WO2019222988A1/zh
Priority to AU2020103572A priority patent/AU2020103572A4/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese

Definitions

  • the invention relates to the field of iron and steel metallurgy, in particular to a superfine-grained high-strength steel plate with a yield strength of 1100 MPa and a manufacturing method thereof.
  • High-strength structural steel plate has very strict requirements on the quality of steel due to its harsh use environment and stress conditions. Its production difficulties are: (1) Ultra-high strength and good plastic toughness are obtained at the same time. Generally, the strength of a steel plate increases and the plastic toughness decreases. How to improve the strength and plastic toughness at the same time is the key to manufacturing high-strength structural steel. (2) How to control the profile of thin gauge steel plates. During the quenching process, due to the superposition of structural stress and thermal stress, it is easy to cause waves at the edges and cores of the steel sheet, which cannot meet the requirements of the shape. (3) How to improve the effective hardening depth of thick gauge steel.
  • Chinese patent CN102747303B describes a high-strength steel plate with a yield strength of 1100 MPa and a method for manufacturing the same.
  • the invention improves the toughness of steel by adding 0.60 to 2.00% Ni, which is very expensive and increases production costs.
  • Chinese patent CN104513936A discloses a quenched and tempered high-strength steel with a yield strength of 1100 MPa and a production method thereof, which also adds 0.30 to 1.50% Ni, and adds a certain amount of Ti and Ca to reduce N by controlling Ti / N and Ca / S.
  • Chinese patent CN106191673A introduces a steel plate with excellent cold bending performance and yield strength greater than 1100 MPa and a preparation method thereof.
  • the straightening process is used to control the shape of the steel plate and the dynamic light reduction process to reduce the central defect of the steel plate.
  • the requirements of the equipment are high, and the promotion and application of this technology are not used.
  • the existing manufacturing technologies that currently involve 1100 MPa ultra-high-strength steel sheets include: (1) improving the toughness of ultra-high-strength steels by adding a large amount of precious metal Ni or Ni + Cu, and the economy is poor; ( 2) Controlling Ti / N and Ca / S by adding Ti and Ca to reduce N and S to improve plasticity and toughness, and the composition design is complicated; (3) Ultra-high martensite strength is obtained by in-line quenching and low-temperature tempering. (4) Control the profile of ultra-high-strength steel through the two straightening processes of warm straightening and strong cold straightening. The process load requires higher equipment.
  • the present invention addresses the above technical problems, overcomes the shortcomings of the prior art, and provides a superfine-grained high-strength steel plate with a yield strength of 1100 MPa.
  • the content of alloying elements is reduced to facilitate the process easily implemented by iron and steel enterprises.
  • the invention reduces a large amount of alloy elements, and does not add precious metal elements such as Ni and Cu, which saves production costs; reduces the carbon equivalent Ceq to less than 0.55%, and improves the welding performance of the steel plate; the yield strength of the steel plate is ⁇ 1100MPa, Has good toughness and plastic deformation ability, especially in the absence of Ni, has very good low temperature toughness.
  • the thickness of the steel plate is 5 to 30 mm.
  • Another object of the present invention is to provide a method for manufacturing a superfine-grained high-strength steel sheet with a yield strength of 1100 MPa, including the following steps:
  • Smelting Smelting by electric furnace or converter according to the above chemical composition, refining by LF furnace, and finally vacuum degassing by VD or RH;
  • the designed liquidus temperature is 1510 °C, and the pouring temperature is 1515 ⁇ 1530 °C. After the slab is taken off line, it enters a heat preservation pit or is covered with a heat preservation cover for 24 hours or more for hydrogen expansion treatment;
  • Heating The slab is heated to 1150 ⁇ 1200 °C, and the temperature is maintained after the temperature of the core of the slab reaches the temperature, and the length of time is 1 ⁇ 1.5mim / mm;
  • the two-stage controlled rolling process is adopted. After the slab is taken out of the furnace, it is depressurized by high-pressure water and then enters the recrystallization zone for rough rolling.
  • the rolling temperature for rough rolling is 1100 ⁇ 1150 °C. %, The thickness of the rough-rolled steel plate to be warmed ⁇ 2.2H, H is the final rolled thickness of the steel plate;
  • the finish rolling rolling temperature is 850 to 950 ° C, the total reduction rate in the finishing rolling stage is ⁇ 70%, and the pass reduction rate is ⁇ 15% , And increase the reduction in the non-recrystallized area;
  • Cooling After rolling, the steel sheet is cooled to 600 ° C at a cooling rate of more than 30 ° C / s and then air-cooled to room temperature;
  • the austenite starting transformation temperature Ac3 of the steel is 814 ° C, and it is rapidly heated to a temperature of 840 to 860 ° C at a heating speed of 50 ° C / min or more. After the furnace temperature reaches the temperature, the holding time is 1 to 1.5min / mm. Machine heat treatment equipment for rapid water cooling to room temperature;
  • Tempering The quenched steel sheet is heated to 200 to 240 ° C, and the heat preservation time is 2 to 3 min / mm after the furnace temperature reaches the temperature, and the air is cooled to room temperature.
  • the aforementioned method for manufacturing a superfine-grained high-strength steel sheet with a yield strength of 1100 MPa adopts electromagnetic stirring during continuous casting, and the electromagnetic stirring parameters are 320A and 6Hz.
  • the above-mentioned process provided by the present invention develops ultra-fine-grained high-strength steel plates with yield strength greater than 1100 MPa and good plasticity and low-temperature toughness through reasonable design of chemical composition and control measures of rolling, cooling process and heat treatment process; Refining grains is the only strengthening method that can improve both strength and plastic toughness.
  • the present invention achieves the toughness requirements of steel plates by refining grains, reducing the addition of alloy elements such as Ni, Mn, and Cr.
  • the carbon equivalent of the steel plate improves the welding performance, on the other hand, it saves costs and improves the competitiveness of the product.
  • the steel plate is heated to the quenching temperature by using a rapid heating process, and the quenching temperature is reduced to 840 to 860 ° C according to the austenite transformation start temperature, thereby avoiding the growth of austenite grains, thereby obtaining an average crystal.
  • the microstructure of the ultrafine-grained high-strength steel plate in the present invention is tempered lath martensite with high dislocation density, the original austenite grain size is ⁇ 10 ⁇ m, the micro-grain size level reaches 11, and the steel plate mechanics Performance meets: tensile strength ⁇ 1250MPa, elongation ⁇ 12%, low-temperature Charpy impact energy at -40 ° C ⁇ 60J, unevenness ⁇ 3mm / m, and good cold bending and welding performance.
  • FIG. 1 is an optical microstructure morphology photograph (500 ⁇ ) of a steel plate in Example 1;
  • FIG. 2 is a scanning electron microscope micrograph of a steel plate in Example 2 (500 ⁇ ); FIG.
  • Example 3 is the morphology and size of the precipitated particles after the ultra-rapid cooling after rolling of the steel sheet in Example 3;
  • Example 4 is a photograph (200 ⁇ ) of a micro grain size of a steel plate in Example 3;
  • FIG. 5 is a photograph (200 ⁇ ) of a micro grain size of a steel sheet in Example 4.
  • FIG. 5 is a photograph (200 ⁇ ) of a micro grain size of a steel sheet in Example 4.
  • C The content of C directly determines the strength level and welding performance of the steel sheet.
  • C is solid-dissolved in the octahedron or tetrahedral gap of the ferrite lattice, causing severe lattice distortion, and interacting with dislocations, which strongly hinders the movement of dislocations, thereby increasing the strength.
  • the C content is too high, the degree of lattice distortion increases, the shear resistance increases, and the plastic toughness of the steel is significantly reduced.
  • the C content is controlled to be 0.15 to 0.22%.
  • Si In steel, Fe atoms in the ferrite crystal lattice are mainly replaced by substitution, which plays a role of solid solution strengthening. In addition, Si can reduce the diffusion ability of C atoms, prevent the formation of carbides during tempering, and thereby improve the tempering resistance of steel. However, when the Si content is too high, the surface quality of the steel sheet is significantly deteriorated, so the Si content is controlled to 0.10 to 0.30%.
  • Mn can suppress the diffusion-type phase transition process, improve the hardenability of the steel, can also reduce the harm of the steel type S, and can also play a solid solution strengthening role.
  • Mn content is controlled to 0.80 to 1.60%.
  • Cr a solid solution strengthening element, which can inhibit the formation of polygonal ferrite and pearlite, promote the transformation of bainite and martensite, and thereby improve the strength.
  • carbides that easily form Cr reduce the toughness of the steel sheet and are not conducive to the welding performance of the steel sheet. Therefore, the Cr content is controlled to 0.20 to 0.70%.
  • Ni It improves the hardenability element and can directly improve the toughness of steel. It is a precious metal and expensive. It is not added in the present invention, which reduces the production cost and improves the competitiveness of the steel type.
  • Mo Improves the hardenability element and promotes the formation of martensite during quenching.
  • Mo can also play a role in refining grains, which is conducive to obtaining a refined martensite structure, but too high Mo content will deteriorate the welding performance of steel, so the Mo content in the present invention is controlled to 0.10 to 0.45%, In order to obtain the matching strength and toughness of the steel plate.
  • Nb In the present invention, the role of Nb is essential. Elemental atoms of Nb and carbides of Nb can significantly pin the austenite grain boundaries and prevent grain growth. In addition, the solute drag of Nb can increase the recrystallization temperature of austenite, which is beneficial to the rolling of the austenite in the non-recrystallized zone, increasing the dislocation density, and providing a refined reheated austenite grain. condition.
  • V is a strong carbide-forming element. These nano-scale V (C, N) are dispersed in the matrix of the steel and can play a role of precipitation strengthening to improve the strength of the steel sheet. In addition, grain boundaries can be pinned during the heating process to prevent austenite grains from growing and play a role in refining the grains. However, when the content of V is too high, the size of the precipitated particles becomes large, which is disadvantageous for toughness. Therefore, the content of V is controlled to 0.020 to 0.060%.
  • Ti has a strong affinity for C and N. Ti and N are liable to form coarse TiN during the solidification of molten steel. TiN is hard and brittle and difficult to deform. As a result, it cannot be deformed cooperatively with the matrix structure during the rolling deformation process, resulting in micro-cracks, which significantly reduces the toughness and cold bending of steel. Performance, therefore, the addition of Ti element is not allowed in the present invention.
  • the most effective element for improving hardenability is to promote the refinement of the martensite structure of the steel sheet during cooling, thereby improving the strength.
  • the B content is controlled to 0.0010 to 0.0030%.
  • Al plays the role of deoxidation in the steel to purify the cleanliness of molten steel; on the other hand, it is also used to fix N in the steel to prevent the formation of coarse TiN to deteriorate the toughness of the steel; it can also protect the B, Avoid the precipitation of BN at the grain boundaries, so as to ensure the role of B to improve the hardenability.
  • the fine AlN can also suppress the growth of austenite grains in the subsequent cooling process and play a role in refining the grains, so the Al content is controlled to 0.02 to 0.06%.
  • N a harmful gas element, forming hard and brittle TiN with Ti significantly deteriorates the toughness of the steel, forming BN with B and enriching it at the grain boundaries, reducing the grain boundary binding energy, and forming AlN with Al can refine the austenite grains, which can Plays the role of fine grain strengthening, so the N content is controlled to ⁇ 0.0040%
  • P segregates at the grain boundaries, which will reduce the binding energy of the grain boundaries. Under the external force of impact, brittle fracture along the crystal is easy to occur, which is the main reason for the first type of tempering brittleness.
  • P and Mn coexist, it will aggravate the tempering brittleness of the steel and significantly deteriorate the toughness of the steel plate.
  • P also deteriorates the welding performance of steel, so the content of P is strictly controlled ⁇ 0.005%.
  • S During the solidification of molten steel, S will segregate, form sulfide inclusions, and reduce the low temperature toughness and cold bending performance of the steel sheet. During the welding process, the formation of sulfides is also prone to thermal cracking, and the SO2 gas formed by the oxidation of S is prone to generate pores in the weld metal, which reduces the performance of the weld joint. Therefore, the content of S is strictly controlled ⁇ 0.002%.
  • H Hazardous gas element. H atoms are easy to accumulate in dislocations in the steel, causing local H partial pressure to be too high, forming micro-cracks to cause H brittleness, reducing the plastic toughness of the steel plate, and seriously endangering the performance of the steel plate. In addition, the excessively high H content is also the main reason for the formation of cutting delay cracks in ultra-high-strength steel plates, so the H content is strictly controlled ⁇ 0.00015%.
  • the production process of the invention is: converter or electric furnace steelmaking ⁇ LF furnace refining ⁇ VD or RH vacuum degassing treatment ⁇ continuous casting ⁇ casting slab hydrogenation treatment ⁇ heating ⁇ rolling ⁇ cooling ⁇ quenching ⁇ tempering.
  • the core of the present invention is to achieve the toughness and plastic deformation ability requirements of ultra-high-strength steel through ultra-refined grains. Therefore, grain refinement measures are carried out in each stage of rolling and heat treatment.
  • the specific process flow as follows:
  • Continuous casting The smelted molten steel is cast into a 150-220mm slab, and the liquidus temperature of the steel is calculated at 1510 ° C based on the chemical composition.
  • the pouring temperature is the liquidus temperature plus 5-20 ° C, which is 1515-1530 ° C.
  • the speed is controlled at 1 ⁇ 1.35m / min.
  • the entire argon gas is used to protect the casting.
  • electromagnetic stirring is used during continuous casting. After the line, enter the heat preservation pit or cover the heat preservation hood for 24 hours or more to carry out the hydrogen expansion treatment.
  • Heating Put the continuous casting slab into a walking-type heating furnace and heat it to 1150 ⁇ 1200 °C. After the temperature of the core of the slab reaches the temperature, heat preservation is started. The heat preservation time is 1 ⁇ 1.5mim / mm to make the chemical in austenite Ingredients are homogenized.
  • the two-stage controlled rolling process is adopted. After the slab is taken out of the furnace, it is depressurized by high-pressure water and then enters the recrystallization zone for rough rolling.
  • the rolling temperature for rough rolling is 1100 ⁇ 1150 °C, and the reduction ratio after three passes is ⁇ 20 % To fully recrystallize and austenite to avoid abnormal coarse grains.
  • the thickness of the rough-rolled steel plate to be warmed is ⁇ 2.2H, where H is the final rolled thickness of the steel plate; the finishing rolling temperature is 850 ⁇ 950 °C.
  • the total rolling reduction in the rolling stage is ⁇ 70%, the pass reduction in passes is ⁇ 15%, and the rolling reduction is increased in the non-recrystallized zone, so that the austenite grains are elongated, and the interfacial area of austenite is fully increased. Reduce the width of austenite grains.
  • Cooling On-line ultra-fast cooling of the rolled steel sheet, cooling to 600 ° C at a cooling rate of 30 ° C / s or higher, and then air-cooling to room temperature, quickly passing through the ferrite transformation region, suppressing (Nb, V) C at The precipitation in high-temperature ferrite causes it to precipitate in the way of uniform nucleation in the subsequent cooling process. At this time, the driving force for precipitation is large, the critical nucleation size is small, and the growth rate of the precipitated particles is slow. The size can be obtained in 10mm.
  • the following (Nb, V) C can play a significant role in pinning and reheating the austenite state and preventing grain growth.
  • the austenite starting transformation temperature Ac3 of the steel measured by the expansion method is 814 ° C.
  • the rolled and cooled steel plate is rapidly heated to a heating rate of 50 ° C / min or more. 840 ⁇ 860 °C, after the furnace temperature reaches the temperature, the holding time is 1 ⁇ 1.5min / mm, and the water is rapidly cooled to room temperature by using the heat treatment equipment of the quenching machine.
  • Tempering The quenched steel sheet is heated to 200-240 ° C, the heat preservation time is 2 to 3 min / mm after the furnace temperature reaches the temperature, and air-cooled to room temperature to eliminate the quenching internal stress of the steel sheet.
  • the ultrafine-grained high-strength steel sheets in Examples 1 to 4 were manufactured according to the chemical components shown in Table 1.
  • the heat-treated steel sheet was subjected to a transverse tensile and a longitudinal impact test.
  • the specific components and process parameters in Examples 1 to 4 are shown in Tables 1 to 3, and the properties of the obtained steel sheet are shown in Table 4.
  • the semi-impact pattern shall be taken in accordance with the requirements of national standards.
  • the microstructure of the steel is a tempered lath martensite structure.
  • the average grain size of the original austenite is less than 10 ⁇ m.
  • the mechanical properties of the steel plate meet: tensile strength ⁇ 1250MPa, elongation ⁇ 12%, low-temperature Charpy impact energy at -40 ° C ⁇ 60J, unevenness ⁇ 3mm / m, and good cold bending and welding performance.
  • the invention optimizes the composition design, controls the rolling and cooling processes, and controls the heat treatment process, and finally realizes the strength of the steel plate while ensuring the plastic toughness through the refined lath martensite structure.
  • the steel plate produced by the present invention has been successfully applied to the equipment of engineering machinery enterprises, replacing imports, and has produced good economic and social benefits.

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Abstract

一种屈服强度1100MPa级超细晶高强钢板，其化学成分及重量百分比如下：C：0.15～0.22％，Si：0.10～0.30％，Mn：0.80～1.60％，Cr：0.20～0.70％，Mo：0.10～0.60％，Nb：0.020～0.050％，V：0.020～0.060％，Ti≤0.008％，B：0.0010～0.0030％，Al：0.02～0.06％，P≤0.005％，S≤0.002％，O≤0.0025％，N≤0.0040％，H≤0.00015％，余量为Fe及不可避免的杂质，碳当量Ceq＝C+Mn/6+(Cr+Mo+V)/5+Ni/15≤0.55％。该钢板成分经济，工艺简单，利用晶粒细化作用使钢板具有优良的综合力学性能，产生良好的经济效益和社会效益。

Description

一种屈服强度1100MPa级超细晶高强钢板及其制造方法 技术领域
本发明涉及钢铁冶金领域，特别是涉及一种屈服强度1100MPa级超细晶高强钢板及其制造方法。
背景技术
近年来，由于受环保问题、能源问题以及成本问题的约束，工程机械和矿上机械等行业向着大型化和轻量化的方向发展，要求工程机械装备要在减轻自重的同时不降低甚至提高载荷能力，因此需要应用更多更高级别的结构钢板。高强度结构钢板主要用于制造起重机的吊臂，工程机械的大梁和自卸车的车体等结构件。我国每年的起重机总产量在1000万吨左右，而吊臂重量占总机重量的15-20％，其每年消耗的结构钢板在150万吨左右。此前，国内各钢厂以生产屈服强度960MPa以下级别的钢板为主，更高级别的结构钢板主要依靠进口。
高强度结构钢板由于其苛刻的使用环境和受力条件，对钢材质量有非常严格的要求。其生产难点在于：(1)超高强度同时获得良好的塑韧性。通常情况，钢板强度增加，塑韧性降低，如何同时提高强度和塑韧性是制造高强度结构钢的关键所在。(2)如何控制薄规格钢板的板型。淬火过程中由于组织应力和热应力的叠加作用，容易使钢板边部和心部产生波浪，无法满足板型要求。(3)如何提高厚规格钢板的有效淬透深度。钢板越厚，表面和心部淬火效果差距增大，容易造成钢板厚度方向性能不均匀。(4)如何改善钢板焊接性能。结构件通常是通过折弯和焊接成型，而通常情况下，通过增加C和合金元素的含量来提高强度会显著的恶化焊接性功能。(5)如何改善钢板冷弯性能。通常情况下，钢板强度越高，塑性形变能力越差，难以满足结构件的折弯需求。(6)如何获得较低的生产成本。通过添加大量的合金元素可以提高强度和塑性，但是过高的生产成本是企业无法接受的。因此必须从组织设计和工艺控制上来实现钢板的强韧化要求。
目前，国内还没有能够批量供应屈服强度1100MPa级高强度结构钢板的企业，主要依靠进口。中国专利CN102747303B介绍了一种屈服强度1100MPa级高强钢板及其制造方法，该发明通过添加0.60～2.00％的Ni来改善钢的韧性，而Ni的价格非常昂贵，增加生产成本。中国专利CN104513936A公布了一种屈服强度1100MPa级调质高强钢及其生产方法，同样添加了0.30～1.50％的Ni，以及添加一定量的Ti和Ca通过控制Ti/N和Ca/S来降低N和S的塑韧性的危害，从而实现高强钢的强韧性匹配，成分设计较复杂。中国专利CN100372962C介绍了一种在线淬火+回火的方法来生产屈服强度1100MPa以上的超高强度钢板，通过添加0.20～1.20％和0～0.5％的Cu来改善韧性，此外，在线淬火的钢板未经过再加热奥氏体化，钢板内应力较大，在后续切割和冷弯过程中容易产生裂纹，实际应用存在一定的局限。中国专利CN106191673A介绍了一种冷弯性能优良屈服强度大于1100MPa的钢板及其制备方法，通过两次矫直工艺来控制钢板的板型和动态轻压下工艺来降低钢板的中心缺陷，工艺复杂对设备的要求较高，不利用该技术的推广应用。
综上所述，目前涉及屈服强度1100MPa级超高强钢板的现有制造技术主要有：(1)通过添加大量的贵重金属Ni或Ni+Cu来改善超高强钢的韧性，经济性较差；(2)通过添加Ti和Ca控制Ti/N和Ca/S来降低N和S来改善塑韧性，成分设计复杂；(3)通过在线淬火和低温回火的方式获得马氏体超高强度。(4)通过温矫和强力冷矫两次矫直工艺来控制超高强钢的板型，工艺负载对设备的要求较高。
发明内容
本发明针对上述技术问题，克服现有技术的缺点，提供一种屈服强度1100MPa级超细晶高强钢板，通过简单的化学成分设计，减少合金元素的含量，以利于钢铁企业容易实现的工艺，同时具有超高强度和良好低温韧性以及塑性变形能力。
为了解决以上技术问题，本发明提供一种屈服强度1100MPa级超细晶高强钢板，其化学成分及重量百分比如下：C：0.15～0.22％，Si：0.10～0.30％，Mn：0.80～1.60％，Cr：0.20～0.70％，Mo：0.10～0.60％，Nb：0.020～0.050％，V：0.020～0.060％，Ti≤0.008％，B：0.0010～0.0030％，Al：0.02～0.06％，P≤0.005％，S≤0.002％，O≤0.0025％，N≤0.0040％，H≤0.00015％，余量为Fe及不可避免的杂质，且碳当量Ceq＝C+Mn/6+(Cr+Mo+V)/5+Ni/15≤0.55％。
技术效果：本发明减少了大量的合金元素，未添加Ni、Cu等贵金属元素，节约了生产成本；将碳当量Ceq降低到0.55％以下，提高了钢板的焊接性能；钢板的屈服强度≥1100MPa，具有良好的强韧性和塑性变形能力，尤其是在未添加Ni的条件下，具有非常好的低温韧性。
本发明进一步限定的技术方案是：
进一步的，钢板的厚度为5～30mm。
本发明的另一目的在于提供一种屈服强度1100MPa级超细晶高强钢板的制造方法，包括以下步骤：
冶炼：按上述化学成分采用电炉或转炉冶炼，经过LF炉精炼，最后经VD或RH真空脱气处理；
连铸：设计液相线温度为1510℃，浇注温度为1515～1530℃，铸坯下线后进入保温坑或加盖保温罩内堆冷24小时以上进行扩氢处理；
加热：将铸坯加热至1150～1200℃，待铸坯心部温度到温后开始保温，保温时长1～1.5mim/mm；
轧制：采用两阶段控轧工艺，首先铸坯出炉后经高压水除磷后进入再结晶区粗轧，粗轧开轧温度1100～1150℃，粗轧后三道次的压下率≥20％，粗轧钢板待温厚度≥2.2H，H为钢板最终轧制厚度；精轧开轧温度为850～950℃,精轧阶 段总压下率≥70％，道次压下率≥15％，并在未再结晶区增大压下量；
冷却：轧制后钢板以30℃/s以上的冷却速度冷却至600℃后再空冷至室温；
淬火：钢的奥氏体开始相变温度Ac3为814℃，以50℃/min以上的加热速度快速加热至840～860℃，炉温到温后保温时长为1～1.5min/mm，利用淬火机热处理装备快速水冷至室温；
回火：将淬火后的钢板加热至200～240℃，炉温到温后保温时长为2～3min/mm，空冷至室温。
前所述的一种屈服强度1100MPa级超细晶高强钢板的制造方法，连铸过程中，全程采用氩气保护浇注。
前所述的一种屈服强度1100MPa级超细晶高强钢板的制造方法，连铸过程中，采用电磁搅拌，电磁搅拌参数为320A，6Hz。
技术效果：本发明提供的上述工艺通过合理设计化学成分并结合轧制、冷却工艺和热处理工艺控制措施，开发出屈服强度大于1100MPa，同时具有良好的塑性和低温韧性的超细晶高强度钢板；细化晶粒是唯一既可以提高强度又可以提高塑韧性的强化方式，本发明通过细化晶粒来实现钢板的韧性要求，减少了Ni、Mn、Cr等合金元素的加入量，一方面降低了钢板的碳当量，提高了焊接性能，另一方面节约了成本，提高了产品的竞争力。
与现有技术相比，本发明的有益效果是：
(1)本发明中通过两阶段控制轧制和轧制后超快冷措施，细化(Nb、V)C的析出粒子尺寸至10mm以下，使其弥散沉淀析出，显著提高其钉扎晶界、阻止晶粒长大的作用；
(2)本发明中通过增加非再结晶区轧制变形量，增大奥氏体晶界面积，减小变形奥氏体的宽度，从而增加再加热奥氏体的形核位置以及奥氏体晶粒横向 生长的阻力，从而起到细化奥氏体晶粒的作用；
(3)本发明中利用快速加热工艺将钢板加热至淬火温度，并依据奥氏体相变开始温度，降低淬火温度至840～860℃，避免奥氏体晶粒的长大，从而得到平均晶粒尺寸小于10μm的奥氏体晶粒；
(4)本发明中采用200～240℃低温回火，消除了淬火内应力，并且保留了高位错密度的板条马氏体，从而获得良好的强韧性；
(5)本发明中超细晶高强钢板的显微组织为高位错密度的回火板条马氏体，原始奥氏体晶粒尺寸＜10μm，显微晶粒度级别达到11级，钢板力学性能满足：抗拉强度≥1250MPa，延伸率≥12％，-40℃时低温夏比冲击功≥60J，不平度≤3mm/m，并且具有良好的冷弯和焊接性能。
附图说明
图1为实施例1中钢板的光学微观组织形貌照片(500×)；
图2为实施例2中钢板的扫描电镜微观组织形貌照片(500×)；
图3为实施例3中钢板轧后超快速冷却后的析出粒子形貌尺寸；
图4为实施例3中钢板的显微晶粒尺寸照片(200×)；
图5为实施例4中钢板的显微晶粒尺寸照片(200×)。
具体实施方式
本实施例提供的一种屈服强度1100MPa级超细晶高强钢板，其化学成分及重量百分比如下：C：0.15～0.22％，Si：0.10～0.30％，Mn：0.80～1.60％，Cr：0.20～0.70％，Mo：0.10～0.60％，Nb：0.020～0.050％，V：0.020～0.060％，Ti≤0.008％，B：0.0010～0.0030％，Al：0.02～0.06％，P≤0.005％，S≤0.002％，O≤0.0025％，N≤0.0040％，H≤0.00015％，余量为Fe及不可避免的杂质，且碳当量Ceq＝C+Mn/6+(Cr+Mo+V)/5+Ni/15≤0.55％。
本发明中的化学成分设计原理：
C：C的含量直接决定了钢板的强度级别和焊接性能。C固溶于铁素体晶格的八面体或四面体间隙中，造成严重的点阵畸变，并与位错发生交互作用，强烈阻碍位错的运动，从而提高强度。但是C含量过高，点阵畸变程度越大，切变抗力增加，显著降低钢的塑韧性，此外在焊接过程中，热影响区容易产生冷裂纹，降低焊接接头的性能，因此本发明中，将C含量控制在0.15～0.22％。
Si：在钢中主要以置换方式替代铁素体晶体点阵中的Fe原子，起到固溶强化的作用。此外，Si可以降低C原子的扩散能力，阻止回火过程中碳化物的形成，从而提高钢的回火抗力。但Si含量过高时，会显著恶化钢板的表面质量，因此将Si含量控制在0.10～0.30％。
Mn：Mn可以抑制扩散型相变过程，提高钢的淬透性，还可以降低钢种S的危害，此外，还能起到固溶强化作用。但Mn含量过高时，容易造成铸坯的中心偏析，恶化钢板内部组织，降低钢板的韧性和焊接性能，因此将Mn含量控制在0.80～1.60％。
Cr：固溶强化元素，可抑制多边形铁素体和珠光体的形成，促进贝氏体和马氏体的转变，从而提高强度。但Cr含量过高时，容易形成Cr的碳化物降低钢板的韧性并且不利于钢板的焊接性能，因此将Cr含量控制在0.20～0.70％。
Ni：提高淬透性元素，并且可以直接提高钢的韧性，其属于贵重金属，价格昂贵，在本发明中不加入该元素，降低生产成本，提高钢种的竞争力。
Mo：提高淬透性元素，促进淬火过程中马氏体的形成。此外，Mo还可以起到细化晶粒作用，有利于得到细化的马氏体组织，但是Mo含量过高会恶化钢的焊接性能，因此本发明中将Mo含量控制在0.10～0.45％，以获得钢板强韧性和焊接性能的匹配。
Nb：在本发明中，Nb的作用至关重要。Nb的单质原子和Nb的碳化物可以显著的钉扎奥氏体晶界，阻止晶粒长大。此外，Nb的溶质拖曳作用可提高奥氏 体的再结晶温度，有利于实现奥氏体的非再结晶区轧制，提高位错密度，为得到细化的的再加热奥氏体晶粒提供条件。
V：V是强碳化物形成元素，这些纳米级的V(C、N)弥散分布在钢的基体中可以起到析出强化的作用，提高钢板的强度。此外，在加热过程中还可以钉扎晶界，阻止奥氏体晶粒长大，起到细化晶粒的作用。但V含量过高时，析出粒子的尺寸变大，反而对韧性不利，因此将V的含量控制在0.020～0.060％。
Ti：Ti与C、N有极强的亲和力。Ti和N在钢液凝固过程中容易形成粗大的TiN，TiN硬而脆，不易变形，导致在轧制变形过程中与基体组织无法协同变形，从而产生微裂纹，显著降低钢的韧性和冷弯性能，因此本发明中不允许加入Ti元素。
B：最有效的提高淬透性元素，促使钢板在冷却过程中得到细化的马氏体组织，从而提高强度。但B含量过高时，会在晶界处富集，降低晶界的结合能，从而降低钢板的韧性，因此将B含量控制在0.0010～0.0030％。
Al：Al一方面在钢中起到脱氧作用，净化钢水洁净度；另一方面还被用来固定钢中的N，避免形成粗大的TiN恶化钢的韧性；还可以起到保护B的作用，避免BN在晶界的析出，从而保证B提高淬透性的作用。此外，细小的AlN在随后的冷却过程中还可抑制奥氏体晶粒的长大，起到细化晶粒的作用，因此将Al含量控制在0.02～0.06％。
N：有害气体元素，与Ti形成硬脆的TiN显著恶化钢的韧性，与B形成BN在晶界处富集，降低晶界结合能，与Al形成AlN可细化奥氏体晶粒，可以起到细晶强化的作用，因此将N含量控制在≤0.0040％
P：P在晶界处偏聚，会降低晶界的结合能，在冲击外力作用下，容易发生沿晶脆性断裂，是造成第一类回火脆性的主要原因。P和Mn共同存在时，会加剧钢的回火脆性，显著恶化钢板的韧性。此外，P还会恶化钢的焊接性能，因此 严格控制P的含量≤0.005％。
S：在钢液凝固过程中，S会发生偏聚，形成硫化物夹杂，降低钢板的低温韧性和冷弯性能。焊接过程中，硫化物的形成还容易产生热裂纹，并且S发生氧化形成的SO2气体在焊缝金属中容易产生气孔，降低焊接接头的性能，因此严格控制S的含量≤0.002％。
O：有害气体元素，含量高，产生的夹杂物多，降低钢板的塑形、韧性和冷弯性能，严格控制O含量≤0.0025％。
H：有害气体元素，H原子在钢中容易在位错处聚集，造成局部H分压过高，形成微裂纹从而产生H脆，降低钢板塑韧性，严重危害钢板的使用性能。此外，H含量过高还是超高强钢板形成切割延迟裂纹的主要原因，因此严格控制H的含量≤0.00015％。
本发明的生产工艺流程为：转炉或电炉炼钢→LF炉精炼→VD或RH真空脱气处理→连铸→铸坯扩氢处理→加热→轧制→冷却→淬火→回火。本发明的核心在于通过超细化的晶粒来实现超高强度钢的强韧性和塑性变形能力要求，因此在轧制和热处理的每个阶段都贯穿着细化晶粒的措施，具体工艺流程如下：
冶炼：按上述化学成分采用150吨转炉或电炉冶炼，经过LF炉精炼，最后经VD或RH真空脱气处理。
连铸：将冶炼的钢水浇铸成150～220mm的铸坯，根据化学成分计算出钢的液相线温度为1510℃，浇注温度为液相线温度加5～20℃即1515～1530℃，拉速控制在1～1.35m/min，为防止钢水氧化，全程氩气保护浇注，为减轻铸坯中心疏松和偏析，在连铸过程中采用电磁搅拌，电磁搅拌参数为320A，6Hz；铸坯下线后进入保温坑或加盖保温罩内堆冷24小时以上进行扩氢处理。
加热：将连铸坯装入步进式加热炉中，加热至1150～1200℃，待铸坯心部温度到温后开始保温，保温时长1～1.5mim/mm，使奥氏体中的化学成分均匀化。
轧制：采用两阶段控轧工艺，首先铸坯出炉后经高压水除磷后进入再结晶区粗轧，粗轧开轧温度1100～1150℃，粗轧后三道次的压下率≥20％，使奥氏体充分再结晶细化，避免出现异常粗大的晶粒，粗轧钢板待温厚度≥2.2H，H为钢板最终轧制厚度；精轧开轧温度为850～950℃,精轧阶段总压下率≥70％，道次压下率≥15％，并在未再结晶区增大压下量，使奥氏体晶粒拉长，充分增加奥氏体的晶界面积，减小奥氏体晶粒的宽度。
冷却：对轧制后的钢板进行在线超快速冷却，以30℃/s以上的冷却速度冷却至600℃后再空冷至室温，快速通过铁素体相变区域，抑制(Nb、V)C在高温铁素体中的析出，使其在后续的冷却过程中以均匀形核的方式沉淀析出，此时析出驱动力大，临界形核尺寸小，析出粒子长大速度缓慢，可得到尺寸在10mm以下的(Nb、V)C，可起到显著钉扎再加热奥氏体境界、阻止晶粒长大的作用。
淬火：根据膨胀法测定钢的奥氏体开始相变温度Ac3为814℃，为了细化再加热奥氏体晶粒，将轧制冷却后的钢板以50℃/min以上的加热速度快速加热至840～860℃，炉温到温后保温时长为1～1.5min/mm，利用淬火机热处理装备快速水冷至室温。
回火：将淬火后的钢板加热至200～240℃，炉温到温后保温时长为2～3min/mm，空冷至室温，消除钢板的淬火内应力。
按表1所示的化学成分，制造实施例1～实施例4中的超细晶高强度钢板。将热处理后的钢板进行横向拉伸、纵向冲击试验。实施例1～实施例4中的具体成分和工艺参数见表1～表3，制得的钢板的各项性能见表4。
表1 实施例1～实施例4中超细晶高强度钢板的化学成分(wt％)
实施例 C Si Mn P S Cr Mo Nb V B Al O N H Ceq
1 0.15 0.30 1.25 0.005 0.001 0.20 0.55 0.04 0.05 0.0020 0.055 0.0012 0.0035 0.0010 0.54
2 0.17 0.25 1.1 0.006 0.001 0.30 0.45 0.04 0.04 0.0018 0.050 0.0015 0.0038 0.0012 0.53
3 0.18 0.20 0.9 0.005 0.001 0.40 0.40 0.03 0.04 0.0016 0.055 0.0010 0.0036 0.0011 0.52
4 0.22 0.15 0.7 0.006 0.001 0.50 0.30 0.03 0.03 0.0015 0.052 0.0013 0.0035 0.0010 0.52
表2 实施例1～实施例4中超细晶高强度钢板的轧制冷却工艺控制
Figure PCTCN2018088336-appb-000001
表3 实施例1～实施例4中超细晶高强度钢板的热处理工艺控制
实施例 厚度/mm 加热速度℃/s 加热温度/℃ 淬火保温时间/min 回火温度/℃ 回火保温时间/℃
1 6 5 840 10 200 20
2 8 5 840 15 210 25
3 20 5 850 25 220 50
4 30 5 860 30 230 75
表4 实施例1～实施例4中超细晶高强度钢板的拉伸、冲击性能
Figure PCTCN2018088336-appb-000002
注：钢板厚度为6mm和8mm时，按国家标准要求取半冲击式样。
如图1-2所示，钢的微观组织为回火板条马氏体组织，如图3-5所示，原始奥氏体平均晶粒尺寸小于10μm。钢板力学性能满足：抗拉强度≥1250MPa，延伸率≥12％，-40℃时低温夏比冲击功≥60J，不平度≤3mm/m，并且具有良好的冷弯和焊接性能。本发明从组织超细化的角度，通过优化成分设计、控制轧制和控制冷却工艺以及控制热处理工艺，最终通过细化的板条马氏体组织来实现钢板强度的同时保证塑韧性。本发明生产的钢板已成功应用于工程机械企业的装备上，替代了进口，产生了良好的经济效益和社会效益。
除上述实施例外，本发明还可以有其他实施方式。凡采用等同替换或等效变换形成的技术方案，均落在本发明要求的保护范围。

Claims (5)

  1. 一种屈服强度1100MPa级超细晶高强钢板，其特征在于，其化学成分及重量百分比如下：C：0.15～0.22％，Si：0.10～0.30％，Mn：0.80～1.60％，Cr：0.20～0.70％，Mo：0.10～0.60％，Nb：0.020～0.050％，V：0.020～0.060％，Ti≤0.008％，B：0.0010～0.0030％，Al：0.02～0.06％，P≤0.005％，S≤0.002％，O≤0.0025％，N≤0.0040％，H≤0.00015％，余量为Fe及不可避免的杂质，且碳当量Ceq＝C+Mn/6+(Cr+Mo+V)/5+Ni/15≤0.55％。
  2. 根据权利要求1所述的一种屈服强度1100MPa级超细晶高强钢板，其特征在于：钢板的厚度为5～30mm。
  3. 如权利要求1～2所述的一种屈服强度1100MPa级超细晶高强钢板的制造方法，其特征在于，包括以下步骤：
    冶炼：按上述化学成分采用电炉或转炉冶炼，经过LF炉精炼，最后经VD或RH真空脱气处理；
    连铸：设计液相线温度为1510℃，浇注温度为1515～1530℃，铸坯下线后进入保温坑或加盖保温罩内堆冷24小时以上进行扩氢处理；
    加热：将铸坯加热至1150～1200℃，待铸坯心部温度到温后开始保温，保温时长1～1.5mim/mm；
    轧制：采用两阶段控轧工艺，首先铸坯出炉后经高压水除磷后进入再结晶区粗轧，粗轧开轧温度1100～1150℃，粗轧后三道次的压下率≥20％，粗轧钢板待温厚度≥2.2H，H为钢板最终轧制厚度；精轧开轧温度为850～950℃,精轧阶段总压下率≥70％，道次压下率≥15％，并在未再结晶区增大压下量；
    冷却：轧制后钢板以30℃/s以上的冷却速度冷却至600℃后再空冷至室温；
    淬火：钢的奥氏体开始相变温度Ac3为814℃，以50℃/min以上的加热速度快速加热至840～860℃，炉温到温后保温时长为1～1.5min/mm，利用 淬火机热处理装备快速水冷至室温；
    回火：将淬火后的钢板加热至200～240℃，炉温到温后保温时长为2～3min/mm，空冷至室温。
  4. 根据权利要求3所述的一种屈服强度1100MPa级超细晶高强钢板的制造方法，其特征在于：所述连铸过程中，全程采用氩气保护浇注。
  5. 根据权利要求3所述的一种屈服强度1100MPa级超细晶高强钢板的制造方法，其特征在于：所述连铸过程中，采用电磁搅拌，电磁搅拌参数为320A，6Hz。
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