WO2019218135A1 - 屈服强度1000MPa级低屈强比超高强钢及其制备方法 - Google Patents

屈服强度1000MPa级低屈强比超高强钢及其制备方法 Download PDF

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WO2019218135A1
WO2019218135A1 PCT/CN2018/086824 CN2018086824W WO2019218135A1 WO 2019218135 A1 WO2019218135 A1 WO 2019218135A1 CN 2018086824 W CN2018086824 W CN 2018086824W WO 2019218135 A1 WO2019218135 A1 WO 2019218135A1
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yield
strength
steel
low
rolling
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French (fr)
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胡军
刘悦
张彬
宋娜
谢辉
杜林秀
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东北大学
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    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/58—Ferrous alloys, e.g. steel alloys containing chromium with nickel with more than 1.5% by weight of manganese
    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • 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
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/06—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
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • 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
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention belongs to the technical field of metallurgical materials, and particularly relates to a low-yield ratio ultra-high strength steel with a yield strength of 1000 MPa and a preparation method thereof.
  • the off-line quenching process can eliminate the post-rolling band structure and the internal stress of the steel plate, and at the same time refine the size of the martensite lath bundle to obtain a fine lath martensite structure with high dislocation density to ensure the steel plate has an ultra-high yield. strength.
  • the invention adopts the composition design of low-carbon Si-Mn-Cr system, and adopts the method of reducing the carbon content and compounding the alloying elements such as Si, Mn and Cr, and adopts the off-line quenching + low temperature tempering treatment to successfully develop a yield strength higher than that.
  • the object of the present invention is to provide a low-yield ratio ultra-high strength steel with a yield strength of 1000 MPa and a preparation method thereof.
  • the composition design and off-line quenching + low temperature tempering treatment process achieve excellent performance and make the process stable and easy to realize. Industrial production.
  • the composition of the present invention has a yield strength of 1000 MPa and a low yield ratio ultrahigh strength steel: C 0.06 to 0.14%, Mn 2.5 to 3.5%, Si 1.0 to 1.6%, Cr 0.5 to 1.2%, and Mo 0.1 to 0.3%. Ni 0.1 ⁇ 0.5%, S 0.002 ⁇ 0.005%, P 0.003 ⁇ 0.010%, Al 0.01 ⁇ 0.05%, N 0.003 ⁇ 0.005%, the balance is Fe and other unavoidable impurities; the yield strength is 1015 ⁇ 1190MPa, The tensile strength is 1290 to 1400 MPa, and the yield ratio is 0.79 to 0.85.
  • the above-mentioned yield strength of 1000MPa grade low yield ratio ultrahigh strength steel elongation is 13.8-16.9%, the impact energy of 10mm thick specimen at -20 °C is ⁇ 100J, and the impact energy of 2.5mm thick specimen at -20°C is ⁇ 36J .
  • microstructure of the above-mentioned low-strength ratio ultra-high-strength steel with a yield strength of 1000 MPa is tempered martensite and retained austenite.
  • the above-mentioned yield strength of 1000 MPa grade low yield ratio ultra high strength steel has a thickness of 3.5 to 35.0 mm.
  • the preparation method of the invention for the yield strength 1000 MPa low yield ratio ultra high strength steel comprises the following steps:
  • the slab after the heat preservation is directly subjected to rough rolling, and then finish rolling to a thickness of 3.5 to 35.0 mm, and then water-cooled to 100 to 500 ° C, and finally air-cooled to room temperature to form a hot-rolled sheet; wherein the final rolling temperature of the rough rolling is 990 ⁇ 1120°C, the rolling temperature of finishing rolling is 900-980° C., and the finishing rolling temperature of finishing rolling is 825-920° C.;
  • the hot-rolled sheet is heated to 850-1000 ° C, and the austenitizing is carried out for 30-90 min, so that the internal structure of the steel sheet is in the austenitizing state; then quenching to ⁇ 100 ° C to ensure that the martensite structure can be obtained;
  • the hot-rolled sheet after quenching is tempered, the tempering temperature is 200-400 ° C, the time is 30-60 min, and then air-cooled to room temperature to obtain a low-yield ratio ultra-high strength steel with a yield strength of 1000 MPa.
  • step 2 the rough rolling is carried out for 1 to 3 passes, the pass reduction ratio is 13 to 29%, and the finish rolling is carried out for 5 to 7 passes, and the pass reduction ratio is 21 to 30%.
  • the cooling rate at the time of water cooling is 25 to 45 ° C / s.
  • composition design principle of the present invention is:
  • Carbon acts as a gap atom in steel and can exhibit solid solution strengthening and effectively increase the strength of the steel sheet.
  • a higher C content increases the carbon equivalent and weld crack sensitivity index, and deteriorates the performance of the weld heat affected zone.
  • the low carbon is used to synergize between the alloy elements to improve the strength; therefore, the C content of the present invention ranges from 0.06 to 0.14%;
  • Manganese and Chromium Manganese is the most basic element in steel, while Mn and Cr are among the most important elements in the present invention. As an austenite stabilizing element, the hardenability of steel can be improved, not only the formation of pro-eutectoid ferrite can be avoided, but also the volume fraction of retained austenite can be increased; therefore, the Mn content in the present invention ranges from 2.5 to 3.5%. , Cr content ranges from 0.5 to 1.2%;
  • Silicon is also one of the most important elements in the present invention.
  • the addition of Si has an important influence on the TRIP effect, which can effectively inhibit the precipitation of cementite during low temperature tempering, and ensure the stability of retained austenite, Si content. Not less than 1.0%, otherwise it is difficult to suppress cementite precipitation, but too high Si content will deteriorate the toughness of martensitic high-strength steel, reduce the weldability of the steel sheet, and affect the surface quality of the steel sheet; therefore, the Si content range of the present invention 1.0 to 1.6%;
  • Molybdenum 0.1% or more of Mo element can improve the hardenability of steel, which is beneficial to form a full martensite structure during quenching; Mo can reduce the temper brittleness of martensite structure, and too high Mo content will lead to an increase in carbon equivalent, Conducive to welding, and increase the cost of the alloy; therefore, the Mo content of the present invention ranges from 0.1 to 0.3%;
  • Nickel Ni element has refined martensite structure to improve the toughness of steel, but too high Ni content will greatly increase the cost of the alloy; therefore, the Ni content of the present invention ranges from 0.1 to 0.5%;
  • S is an impurity element in steel.
  • the affinity of S and Mn is strong, and it is easy to form MnS.
  • the transverse tensile properties of the product are reduced during the rolling process of the steel sheet.
  • the excessive S content is easy to cause hot brittleness, and should be controlled as much as possible.
  • the content is to a lower level; therefore, the S content of the present invention ranges from 0.002 to 0.005%;
  • P is also an impurity element in steel. Although an appropriate amount of P element is beneficial to prevent the precipitation of cementite and retain more metastable austenite, too high P content causes cold brittleness, and for ultra high strength steel. Plasticity is unfavorable; therefore, the P content of the present invention ranges from 0.003 to 0.010%;
  • Al can also inhibit the precipitation of cementite, and combine with N to form AlN, it can effectively refine the grains and improve the impact toughness.
  • the excessive Al content leads to the viscous molten steel in the continuous casting process, which reduces the efficiency of casting.
  • the content of Al exceeding 0.06% is liable to cause oxide inclusion defects of Al; therefore, the Al content of the present invention ranges from 0.01 to 0.05%;
  • N is also an inevitable impurity element in steel, and its content should be controlled as low as possible; therefore, the N content of the present invention ranges from 0.003 to 0.005%.
  • the quenching process can obtain a fine lath martensite structure with high dislocation density, which greatly increases the yield strength of the steel plate; in the low-temperature tempering process, the carbon atom is effective by the diffusion partitioning of C element. Enriched into austenite, the addition of Si prevents the precipitation of carbides, so that the steel plate obtains a small amount of retained austenite structure; at the same time, the tempering process reduces the dislocation density, reduces the residual internal stress of the steel plate after quenching, and removes the hydrogen of the steel plate. , thereby obtaining an ultra-high strength steel plate having a yield strength of 1000 MPa and a low yield ratio;
  • off-line quenching can precisely control the temperature at which quenching starts; at the same time, the off-line quenching process uses hot-air cooling after hot rolling, which is easier to control the flatness of the plate than the on-line quenching, and the quenching machine has strong quenching ability.
  • the production thickness range is relatively wide; and with the subsequent tempering process, the microstructure and properties of the steel are further optimized;
  • 1 is a schematic flow chart showing a preparation method of a low-yield ratio ultra-high strength steel with a yield strength of 1000 MPa;
  • Example 2 is a SEM organization diagram of a low-strength ratio ultra-high strength steel of a yield strength of 1000 MPa in Example 1;
  • Example 3 is a TEM structure diagram of the low-strength ratio ultra-high strength steel of the yield strength of 1000 MPa in Example 2.
  • test methods described in the examples are conventional methods unless otherwise specified; the reagents and materials are commercially available unless otherwise specified.
  • the hot rolling mill used in the embodiment is the ⁇ 450mm hot rolling mill of the State Key Laboratory of Rolling Technology and Continuous Rolling Automation of Northeastern University.
  • the heating furnace used in the low temperature tempering in the embodiment is a high temperature box type electric resistance furnace, and the model number is RX-36-10.
  • the scanning electron microscope used in the examples was a Zeiss Ultra 55 scanning electron microscope.
  • the transmission electron microscope used in the examples was a FEI Tecnai G 2 F20 field emission transmission electron microscope.
  • the slab in the embodiment is a slab made by smelting, casting and forging.
  • the thickness of the slab 140mm As a heating furnace, heat 4h at 1200 deg.] C; wherein the steel slab composition in percentage by weight C 0.06%, Mn 3.5%, Si 1.6%, Cr 1.2%, Mo 0.3%, Ni 0.5%, S 0.002%, P 0.010%, Al 0.05%, N 0.005%, the balance is Fe and other unavoidable impurities;
  • the slab after the heat preservation is rough-rolled and finish-rolled to a thickness of 35.0 mm, then water-cooled to 500 ° C, and the cooling rate is 25 ° C / s; finally, air-cooled to room temperature to make a hot-rolled sheet; wherein the final rolling temperature of the rough rolling is 1120 ° C , rough rolling for 3 passes, rough rolling to thickness of 90mm, pass reduction rate of 13 to 29%; finish rolling rolling temperature of 980 ° C, finish rolling finishing temperature of 920 ° C, finishing rolling for 7 Second, the pass reduction rate is 21 to 30%;
  • the hot-rolled sheet is heated to 1000 ° C, and austenitized for 90 min to make the internal structure of the steel sheet in an austenitizing state; then quenched to 100 ° C to ensure that martensite structure can be obtained;
  • the quenched hot-rolled sheet is tempered, tempered at a temperature of 400 ° C for 60 min, and then air-cooled to room temperature to obtain a low-yield ratio ultra-high strength steel having a yield strength of 1000 MPa;
  • the yield strength of the 1000MPa grade low yield ratio ultra high strength steel is 1015MPa
  • the tensile strength is 1290MPa
  • the yield ratio is 0.79
  • the elongation is 16.9%
  • the impact energy of the 10mm thick specimen at -20°C is 138J.
  • the microstructure is tempered martensite and retained austenite.
  • the SEM structure is shown in Figure 2.
  • the composition of the billet is: C 0.10%, Mn 3.0%, Si 1.3%, Cr 0.8%, Mo 0.2%, Ni 0.3%, S 0.008%, P 0.003%, Al 0.01%, N 0.004%
  • the balance is Fe and other unavoidable impurities; the thickness of the billet is 80 mm; the temperature is kept at 1150 ° C for 3 h;
  • Ultra high strength steel has a yield strength of 1083MPa, tensile strength of 1340MPa, yield ratio of 0.80, elongation of 14.7%, and impact energy of 10mm thick specimen at -20°C 109J
  • the TEM organization is shown in Figure 3.
  • the composition of the billet is by weight: C 0.14%, Mn 2.5%, Si 1.0%, Cr 0.5%, Mo 0.1%, Ni 0.1%, S 0.005%, P 0.010%, Al 0.03%, N 0.003%
  • the balance is Fe and other unavoidable impurities; the thickness of the billet is 50 mm; the temperature is kept at 1120 ° C for 2 h;
  • Yield strength 1000MPa grade low yield ratio ultra high strength steel yield strength is 1190MPa
  • tensile strength is 1400MPa
  • yield ratio is 0.85
  • elongation is 13.8%

Abstract

一种屈服强度1000MPa级低屈强比超高强钢及其制备方法，超高强钢的成分按重量百分比为：C 0.06～0.14％，Mn 2.5～3.5％，Si 1.0～1.6％，Cr 0.5～1.2％，Mo 0.1～0.3％，Ni 0.1～0.5％，S 0.002～0.005％，P 0.003～0.010％，Al 0.01～0.05％，N 0.003～0.005％，余量为Fe，屈服强度1015～1190MPa，屈强比0.79～0.85；方法为：(1)将钢坯在1120～1200℃保温；(2)进行粗轧和精轧，然后水冷后再空冷；(3)加热至850～1000℃进行奥氏体化，然后淬火；(4)在200～400℃回火处理，随后空冷。

Description

屈服强度1000MPa级低屈强比超高强钢及其制备方法 技术领域
本发明属于冶金材料技术领域，特别涉及一种屈服强度1000MPa级低屈强比超高强钢及其制备方法。
背景技术
随着国内外钢铁行业竞争的不断加剧，各钢铁企业着力于开发具有高经济附加值的产品。目前，铁路交通、工程机械、重载汽车等领域的飞速发展，给钢铁材料的综合力学性能提出了新的要求，低强度钢板已经难以满足下游客户的使用需求，因此这促使钢材的强度级别不断攀升，从Q500MPa级别快速升高至Q550MPa、Q620MPa以及Q690MPa级别，并向更高级别钢板提出了明确的需求。由于采用超高强度薄钢板替代低强度厚钢板能够降低设备自重并提高装备安全性，可实现降低成本、轻量化、节能降耗的目的，因此屈服强度1000MPa级超高强钢的推广应用具有广阔的市场前景。
传统高强钢通过固溶强化、析出强化或者细化晶粒的方式来提高强度，而超高强钢则更多依赖组织强化的方式来达到超高强度。虽然目前TMCP工艺已经广泛的应用于高强钢制备过程，但对于超高强钢难以实现屈服强度1000MPa并满足组织性能均匀性。将经过TMCP后的钢材离线再加热到奥氏体化温度，保温一段时间后淬火到室温，随后再进行低温回火热处理工艺，可以避免出现直接通过TMCP工艺造成的内部组织不均匀和综合性能波动大的问题。此外，离线淬火过程可以消除轧后带状组织以及钢板内应力，同时细化马氏体板条束尺寸，从而获得高位错密度的精细板条马氏体组织，来保证钢板具有超高的屈服强度。
国内诸多钢铁企业已经相继开发出不同强度级别的超高强钢板，但由于受到生产设备、工艺流程、合金成本等因素的限制，使得在产品规格和性能等方面仍难以满足高端市场需求。本发明采用低碳Si-Mn-Cr系的成分设计，通过降低碳含量，复合添加Si、Mn、Cr等合金元素的方法，采用离线淬火+低温回火处理成功开发出一种屈服强度高于1000MPa级别的低屈强比超高强钢板。
发明内容
本发明的目的是提供一种屈服强度1000MPa级低屈强比超高强钢及其制备方法，通过成分设计以及离线淬火+低温回火处理工艺，达到优良性能的同时，使工序过程稳定，容易实现工业化生产。
本发明的屈服强度1000MPa级低屈强比超高强钢的成分按重量百分比为：C 0.06～0.14％，Mn 2.5～3.5％，Si 1.0～1.6％，Cr 0.5～1.2％，Mo 0.1～0.3％，Ni 0.1～0.5％，S 0.002～0.005％，P  0.003～0.010％，Al 0.01～0.05％，N 0.003～0.005％，余量为Fe和其他不可避免的杂质；其屈服强度为1015～1190MPa，抗拉强度为1290～1400MPa，屈强比为0.79～0.85。
上述的屈服强度1000MPa级低屈强比超高强钢的延伸率为13.8～16.9％，-20℃下10mm厚试样的冲击功≥100J，-20℃下2.5mm厚试样的冲击功≥36J。
上述的屈服强度1000MPa级低屈强比超高强钢的显微组织为回火马氏体和残余奥氏体。
上述的屈服强度1000MPa级低屈强比超高强钢的厚度为3.5～35.0mm。
本发明的屈服强度1000MPa级低屈强比超高强钢的制备方法包括以下步骤：
1、将厚度为50～140mm的钢坯至于加热炉中，在1120～1200℃保温2～4h；其中钢坯的 成分按重量百分比为：C 0.06～0.14％，Mn 2.5～3.5％，Si 1.0～1.6％，Cr 0.5～1.2％，Mo 0.1～0.3％，Ni 0.1～0.5％，S 0.002～0.005％，P 0.003～0.010％，Al 0.01～0.05％，N 0.003～0.005％，余量为Fe和其他不可避免的杂质；；
2、将保温后的钢坯直接进行粗轧，再进行精轧至厚度3.5～35.0mm，然后水冷至100～500℃，最后空冷至室温制成热轧板；其中粗轧的终轧温度为990～1120℃，精轧的开轧温度为900～980℃，精轧的终轧温度为825～920℃；
3、将热轧板加热至850～1000℃，保温30～90min进行奥氏体化，使其钢板内部组织处于奥氏体化状态；然后淬火至≤100℃，确保能够获得马氏体组织；
4、将淬火后的热轧板进行回火处理，回火温度200～400℃，时间30～60min，随后空冷至室温，获得屈服强度1000MPa级低屈强比超高强钢。
上述的步骤2中，粗轧进行1～3道次，道次压下率为13～29％；精轧进行5～7道次，道次压下率为21～30％。
上述的步骤2中，水冷时的冷却速率为25～45℃/s。
本发明的成分设计原理为：
碳：碳作为钢中的间隙原子，能够发挥固溶强化作用，有效的提高钢板的强度，但是较高的C含量会提高碳当量和焊接裂纹敏感性指数，恶化焊接热影响区性能。根据“多元少量”的成分设计思路，采用低碳配合各合金元素之间的协同作用来提高强度；因此，本发明的C含量范围为0.06～0.14％；
锰和铬：锰是钢中最基本的元素，同时Mn和Cr是本发明中最重要的元素之一。作为奥氏体稳定化元素，可以提高钢的淬透性，不但能够避免先共析铁素体形成，而且可以提高残余奥氏体体积分数；因此，本发明中Mn含量范围为2.5～3.5％，Cr含量范围为0.5～1.2％；
硅：硅也是本发明中最重要的元素之一，Si的添加对TRIP效应具有重要影响，能够在 低温回火过程中有效抑制渗碳体的析出，保证残余奥氏体的稳定性，Si含量不能低于1.0％，否则难以抑制渗碳体析出，但是过高的Si含量会恶化马氏体高强钢韧性，降低钢板的焊接性能，并影响钢板的表面质量；因此，本发明的Si含量范围为1.0～1.6％；
钼：0.1％以上的Mo元素可以提高钢的淬透性，有利于淬火时形成全马氏体组织；Mo可降低马氏体组织的回火脆性，Mo含量太高会导致碳当量提高，不利于焊接，而且提高合金成本；因此，本发明Mo含量范围为0.1～0.3％；
镍：Ni元素具有细化马氏体组织，改善钢材强韧性的作用，但是过高的Ni含量会大幅提高合金成本；因此，本发明Ni含量范围为0.1～0.5％；
硫：S是钢中的杂质元素，S和Mn的亲和力较强，易于形成MnS，在钢板轧制过程中降低产品的横向拉伸性能，过高的S含量易引起热脆性，应尽量控制其含量至较低水平；因此，本发明S含量的范围为0.002～0.005％；
磷：P也是钢中的杂质元素，虽然适量的P元素有利于阻止渗碳体的析出，保留更多的亚稳态奥氏体，但过高的P含量造成冷脆性，对超高强钢的塑性是不利的；因此，本发明P含量的范围为0.003～0.010％；
铝：虽然Al也能抑制渗碳体的析出，与N结合形成AlN，能够有效细化晶粒，提高冲击韧性，但过高的Al含量导致连铸过程中钢液粘稠，降低浇钢效率，同时Al超过0.06％含量容易产生Al的氧化物夹杂缺陷；因此，本发明Al含量的范围为0.01～0.05％；
氮：N也是钢中不可避免的杂质元素，应尽量控制其含量至较低水平；因此，本发明N含量的范围为0.003～0.005％。
与现有技术相比，本发明的优势在于：
(1)采用低碳化学成分设计，淬火过程可以获得高位错密度的精细板条马氏体组织，大幅提高钢板屈服强度；在低温回火过程中，通过C元素的扩散配分使碳原子有效的富集到奥氏体中，Si的添加防止碳化物的析出，使钢板获得少量的残余奥氏体组织；同时回火过程降低位错密度，减少淬火后钢板的残余内应力，排除钢板的氢气，从而获得具有屈服强度1000MPa级低屈强比的超高强钢板；
(2)采用合理的热轧TMCP工艺，与轧后离线淬火+低温回火工艺制度配合，对奥氏体化温度、回火温度和保温时间分别进行控制；相比于热轧后直接进行冷却的钢材，离线淬火能够精准的控制开始淬火的温度；与此同时，离线淬火工艺采用热轧后热矫空冷，与在线淬火相比更容易控制板型平直度，且淬火机淬火能力强，生产厚度范围比较宽；再配合后续的回火工艺，进一步优化钢材的组织与性能；
(3)产品的低温冲击韧性良好，工艺过程简单，无需冷轧亦能实现3.5mm厚高强度薄板的制备，便于实现工业化生产，提高生产效率。
附图说明
图1为本发明的屈服强度1000MPa级低屈强比超高强钢的制备方法流程曲线示意图；
图2为实施例1的屈服强度1000MPa级低屈强比超高强钢的SEM组织图；
图3为实施例2的屈服强度1000MPa级低屈强比超高强钢的TEM组织图。
具体实施方式
下述非限制性实施例可以使本领域的普通技术人员更全面地理解本发明，但不以任何方式限制本发明。
实施例中所述试验方法，如无特殊说明，均为常规方法；所述试剂和材料，如无特殊说明，均可从商业途径获得。
实施例采用的热轧机为东北大学轧制技术及连轧自动化国家重点实验室的Φ450mm热轧机。
实施例中低温回火时采用的加热炉为高温箱式电阻炉，型号为RX-36-10。
实施例中所用扫描电子显微镜为Zeiss Ultra 55扫描电子显微镜。
实施例中所用透射电子显微镜为FEI Tecnai G 2 F20场发射透射电子显微镜。
实施例中的钢坯是经冶炼、浇铸和锻造后制成的钢坯。
实施例1
流程如图1所示；
将厚度为140mm的钢坯至于加热炉中，在1200℃保温4h；其中钢坯的 成分按重量百分比为C 0.06％，Mn 3.5％，Si 1.6％，Cr 1.2％，Mo 0.3％，Ni 0.5％，S 0.002％，P 0.010％，Al 0.05％，N 0.005％，余量为Fe和其他不可避免的杂质；
将保温后的钢坯经粗轧和精轧至厚度35.0mm，再水冷至500℃，冷却速率为25℃/s；最后空冷至室温制成热轧板；其中粗轧的终轧温度为1120℃，粗轧进行3道次，粗轧至厚度90mm，道次压下率为13～29％；精轧的开轧温度为980℃，精轧的终轧温度为920℃，精轧进行7道次，道次压下率为21～30％；
将热轧板加热至1000℃，保温90min进行奥氏体化，使其钢板内部组织处于奥氏体化状态；然后淬火至100℃，确保能够获得马氏体组织；
将淬火后的热轧板进行回火处理，回火温度400℃，时间60min，随后空冷至室温，获得屈服强度1000MPa级低屈强比超高强钢；
屈服强度1000MPa级低屈强比超高强钢的屈服强度为1015MPa，抗拉强度为1290MPa，屈强比为0.79，延伸率为16.9％，-20℃下10mm厚试样的冲击功138J，显微组织为回火马氏体和残余奥氏体，SEM组织如图2所示。
实施例2
方法同实施例1，不同点在于：
(1)钢坯的成分按重量百分比为：C 0.10％，Mn 3.0％，Si 1.3％，Cr 0.8％，Mo 0.2％，Ni 0.3％，S 0.008％，P 0.003％，Al 0.01％，N 0.004％，余量为Fe和其他不可避免的杂质；钢坯的厚度为80mm；在1150℃保温3h；
(2)粗轧和精轧至厚度12.0mm；其中粗轧的终轧温度为1100℃，粗轧进行1道次，粗轧至60mm，压下率为25％；精轧的开轧温度为950℃，精轧的终轧温度为890℃，精轧进行6道次，道次压下率为21～30％；水冷至360℃，冷却速率为35℃/s；
(3)加热至920℃，保温60min；淬火至60℃；
(4)回火温度300℃，时间45min；
(5)屈服强度1000MPa级低屈强比超高强钢的屈服强度为1083MPa，抗拉强度为1340MPa，屈强比为0.80，延伸率为14.7％，-20℃下10mm厚试样的冲击功109J，TEM组织如图3所示。
实施例3
方法同实施例1，不同点在于：
(1)钢坯的成分按重量百分比为：C 0.14％，Mn 2.5％，Si 1.0％，Cr 0.5％，Mo 0.1％，Ni 0.1％，S 0.005％，P 0.010％，Al 0.03％，N 0.003％，余量为Fe和其他不可避免的杂质；钢坯的厚度为50mm；在1120℃保温2h；
(2)粗轧和精轧至厚度3.5mm；其中粗轧的终轧温度为990℃，粗轧进行3道次，道次压下率为13～29％，粗轧至20mm；精轧的开轧温度为900℃，精轧的终轧温度为825℃，精轧进行5道次，道次压下率为21～30％；水冷至100℃，冷却速率为45℃/s；
(3)加热至850℃，保温30min；淬火至20℃；
(4)回火温度200℃，时间30min；
(5)屈服强度1000MPa级低屈强比超高强钢的屈服强度为1190MPa，抗拉强度为1400MPa，屈强比为0.85，延伸率为13.8％，-20℃下2.5mm厚试样的冲击功36J。

Claims (7)

  1. 一种屈服强度1000MPa级低屈强比超高强钢，其特征在于成分按重量百分比为：C 0.06～0.14％，Mn 2.5～3.5％，Si 1.0～1.6％，Cr 0.5～1.2％，Mo 0.1～0.3％，Ni 0.1～0.5％，S 0.002～0.005％，P 0.003～0.010％，Al 0.01～0.05％，N 0.003～0.005％，余量为Fe和其他不可避免的杂质；其屈服强度为1015～1190MPa，抗拉强度为1290～1400MPa，屈强比为0.79～0.85。
  2. 根据权利要求1所述的屈服强度1000MPa级低屈强比超高强钢，其特征在于其延伸率为13.8～16.9％，-20℃下10mm厚试样的冲击功≥100J，-20℃下2.5mm厚试样的冲击功≥36J。
  3. 根据权利要求1所述的屈服强度1000MPa级低屈强比超高强钢，其特征在于其显微组织为回火马氏体和残余奥氏体。
  4. 根据权利要求1所述的屈服强度1000MPa级低屈强比超高强钢，其特征在于其厚度为3.5～35.0mm。
  5. 一种权利要求1所述的屈服强度1000MPa级低屈强比超高强钢的制备方法，其特征在于包括以下步骤：
    (1)将厚度为50～140mm的钢坯至于加热炉中，在1120～1200℃保温2～4h；其中钢坯的成分按重量百分比为：C 0.06～0.14％，Mn 2.5～3.5％，Si 1.0～1.6％，Cr 0.5～1.2％，Mo 0.1～0.3％，Ni 0.1～0.5％，S 0.002～0.005％，P 0.003～0.010％，Al 0.01～0.05％，N 0.003～0.005％，余量为Fe和其他不可避免的杂质；
    (2)将保温后的钢坯直接进行粗轧，再进行精轧至厚度3.5～35.0mm，然后水冷至100～500℃，最后空冷至室温制成热轧板；其中粗轧的终轧温度为990～1120℃，精轧的开轧温度为900～980℃，精轧的终轧温度为825～920℃；
    (3)将热轧板加热至850～1000℃，保温30～90min进行奥氏体化，使其钢板内部组织处于奥氏体化状态；然后淬火至≤100℃，确保能够获得马氏体组织；
    (4)将淬火后的热轧板进行回火处理，回火温度200～400℃，时间30～60min，随后空冷至室温，获得屈服强度1000MPa级低屈强比超高强钢。
  6. 根据权利要求5所述的屈服强度1000MPa级低屈强比超高强钢的制备方法，其特征在于步骤(2)中，粗轧进行1～3道次，道次压下率为13～29％；精轧进行5～7道次，道次压下率为21～30％。
  7. 根据权利要求5所述的屈服强度1000MPa级低屈强比超高强钢的制备方法，其特征在于步骤(2)中，水冷时的冷却速率为25～45℃/s。
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