WO2021238916A1 - 一种超高强双相钢及其制造方法 - Google Patents

一种超高强双相钢及其制造方法 Download PDF

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WO2021238916A1
WO2021238916A1 PCT/CN2021/095807 CN2021095807W WO2021238916A1 WO 2021238916 A1 WO2021238916 A1 WO 2021238916A1 CN 2021095807 W CN2021095807 W CN 2021095807W WO 2021238916 A1 WO2021238916 A1 WO 2021238916A1
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phase steel
strength
ultra
dual
temperature
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PCT/CN2021/095807
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French (fr)
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李伟
朱晓东
薛鹏
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宝山钢铁股份有限公司
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Priority to CA3180467A priority Critical patent/CA3180467A1/en
Priority to EP21813825.3A priority patent/EP4159886A4/en
Priority to US17/927,781 priority patent/US20230227930A1/en
Priority to JP2022572701A priority patent/JP2023527389A/ja
Publication of WO2021238916A1 publication Critical patent/WO2021238916A1/zh

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    • 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
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
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    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
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    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
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    • 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
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    • C25—ELECTROLYTIC OR ELECTROPHORETIC PROCESSES; APPARATUS THEREFOR
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    • C25D3/00—Electroplating: Baths therefor
    • C25D3/02—Electroplating: Baths therefor from solutions
    • C25D3/22—Electroplating: Baths therefor from solutions of zinc
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    • C25D5/00—Electroplating characterised by the process; Pretreatment or after-treatment of workpieces
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    • C21D2211/00—Microstructure comprising significant phases
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    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Definitions

  • the invention relates to a metal material and a manufacturing method thereof, in particular to a dual-phase steel and a manufacturing method thereof.
  • the demand for strength grades on the market is mainly 80 kg and 100 kg grades.
  • the highest strength grade is currently 1180DP, with a tensile strength greater than or equal to 1200 MPa, a yield strength of about 850 MPa, and a total elongation of about 10%.
  • the production of cold-rolled dual-phase steel adopts the continuous annealing process in the critical zone, and its tensile strength is determined by the martensite fraction in the annealed structure. The higher the martensite fraction, the higher the tensile strength, which requires it during production.
  • the highest strength grade of dual-phase steel that can be produced commercially is 1180MPa, that is, DP 1180 steel.
  • the publication number is CN109504930A, and the publication date is March 22, 2019.
  • the Chinese patent document entitled "Hot-dip galvanized steel sheet with tensile strength greater than 1300MPa and its production method” discloses a hot-dip galvanized steel sheet with tensile strength greater than 1300MPa Zinc steel plate and its production method.
  • the chemical composition of the hot-dip galvanized steel plate substrate and its mass percentage content are: C: 0.1-0.2%, Mn: 1.3-2.0%, S ⁇ 0.005%, P ⁇ 0.02%, Si : 0.2 ⁇ 0.3%, Als: 0.4 ⁇ 1.0%, Nb: 0.01 ⁇ 0.03%, Ti: 0.04 ⁇ 0.08%, B: 0.001 ⁇ 0.004%, Mo: 0.2 ⁇ 0.3%, Cr: 0.05 ⁇ 0.10%, V: 0.01 to 0.02%, the balance is Fe and unavoidable impurities.
  • the heating temperature is 1200 ⁇ 1320°C, and the heating time is 120 ⁇ 200min;
  • the hot rolling process rough rolling is 3 ⁇ 7 passes;
  • the finishing rolling inlet temperature is 1020 ⁇ 1080°C, and the final rolling temperature is 820 ⁇ 880°C;
  • the production method includes slab heating, hot rolling, pickling, continuous hot-dip galvanizing, smoothing and passivation processes;
  • the continuous hot-dip galvanizing process the soaking temperature is 760 ⁇ 840°C, holding time 100 ⁇ 200s, slow cooling temperature 680 ⁇ 740°C, slow cooling cooling rate 10 ⁇ 20°C/s, fast cooling temperature 420 ⁇ 450°C, fast cooling cooling rate 35 ⁇ 65°C/s, plating
  • the zinc temperature is 458 ⁇ 462°C, and the galvanizing time is 5 ⁇ 15s.
  • the publication number is CN108486494A, the publication date is September 4, 2018, and the Chinese patent document entitled "Vanadium microalloyed 1300MPa grade high-strength hot-rolled steel sheet and cold-rolled dual-phase steel sheet production method" discloses a vanadium microalloy
  • the chemical composition of 1300MPa grade high-strength hot-rolled steel sheet and cold-rolled dual-phase steel sheet is as follows: 0.10-0.30wt% C, 1.50-4.50wt% Mn, 0.00-0.120wt% Al, 0.00-0.90wt% Si, 0.05-0.50%V, P ⁇ 0.020wt%, S ⁇ 0.02wt%, Fe: balance.
  • the high-strength steel pass combines the precipitation strengthening of nano vanadium carbide particles with martensitic transformation strengthening, which significantly improves the strength of the existing dual-phase steel while ensuring higher production efficiency.
  • the publication number is CN109628846A, the publication date is April 16, 2019, and the Chinese patent document titled "1300MPa-grade ultra-high-strength cold-rolled steel sheet for automobiles and its production method" discloses a hot-formed steel sheet and a manufacturing method.
  • the chemical composition is: C: 0.1 ⁇ 0.2%, Mn: 1.3 ⁇ 2.0%, S ⁇ 0.005%, P ⁇ 0.02%, Si: 0.2 ⁇ 0.3%, Als: 0.4 ⁇ 1.0%, Nb: 0.01 ⁇ 0.03%, Ti : 0.04 ⁇ 0.08%, B: 0.001 ⁇ 0.004%, Mo: 0.2 ⁇ 0.3%, Cr: 0.05 ⁇ 0.10%, V: 0.01 ⁇ 0.02%, Fe: balance.
  • the production method includes the processes of steelmaking, continuous casting, hot rolling, pickling, continuous annealing, and leveling and straightening; in the hot rolling process, the slab heating temperature is ⁇ 1200°C, the rough rolling is 3 to 7 passes, and the After rolling, the thickness of the intermediate billet is 28-40mm, the finishing rolling inlet temperature is 1020-1100°C, the final rolling temperature is 820-900°C, and the coiling temperature is 550-650°C; in the pickling process, cold rolling and cold rolling are performed after pickling.
  • the heat preservation temperature of the soaking section is 760 ⁇ 840°C, and the heat preservation time is 60 ⁇ 225s; the heat preservation temperature of the overaging section is 250 ⁇ 320°C, and the heat preservation time of the overaging section is 300 ⁇ 1225s. .
  • One of the objectives of the present invention is to provide an ultra-high-strength dual-phase steel which adopts a reasonable chemical element composition design and adopts a medium-Si and low-Al design to reduce the use of alloy elements such as Si and Al, and avoid causes Surface quality caused by high Si and slab defects caused by high Al.
  • the ultra-high-strength dual-phase steel of the present invention does not use precious alloy elements such as Cr and Mo, which effectively controls the alloy cost, while reducing the content of impurity elements P and S, which is beneficial to the improvement of performance and the improvement of delayed cracking.
  • the yield strength of the ultra-high-strength dual-phase steel is ⁇ 900MPa, preferably ⁇ 930MPa, the tensile strength is ⁇ 1300MPa, preferably ⁇ 1320MPa, the elongation after fracture is ⁇ 5%, preferably ⁇ 5.5%, and the initial hydrogen content is ⁇ 10ppm, preferably ⁇ 7ppm ;
  • the preset stress is greater than or equal to one time the tensile strength, the delayed cracking will not occur when immersed in 1mol/L hydrochloric acid for 300 hours.
  • the preset stress is 1.2 times the tensile strength
  • 1mol/L Hydrochloric acid immersion for 300 hours does not cause delayed cracking, which can be effectively applied to the manufacture of automobile safety structural parts, and has good promotion and application value and prospects.
  • the present invention provides an ultra-high-strength dual-phase steel, the matrix structure of which is ferrite + martensite, in which ferrite and martensite are uniformly distributed in an island shape.
  • the ultra-high-strength dual-phase steel In addition to Fe, it also contains the following chemical elements with the following mass percentages:
  • the mass percentage of each chemical element is:
  • C In the ultra-high-strength dual-phase steel of the present invention, C is a solid solution strengthening element, which is a guarantee for the material to obtain high strength. However, it should be noted that the higher the C content in the steel, the harder the martensite and the greater the tendency for delayed cracking to occur. Therefore, when designing the product, try to choose a low-carbon design, and control the mass percentage of C in the ultra-high-strength dual-phase steel of the present invention to be between 0.12 and 0.2%.
  • the mass percentage of C can be controlled between 0.14-0.18%.
  • Si plays a role in increasing the elongation in the steel. Si also has a great influence on the structure of steel, promoting the purification of ferrite and the formation of retained austenite. At the same time, it can improve the tempering resistance of martensite and inhibit the precipitation and growth of Fe 3 C, so that the precipitates formed during tempering are mainly epsilon carbides. But it should be noted that when the mass percentage of Si in the steel is less than 0.5%, it will affect the elongation and tempering resistance of the steel, and if the mass percentage of Si is higher than 1.0%, it will bring other metallurgical quality defects. . Therefore, the mass percentage of Si in the ultra-high-strength dual-phase steel of the present invention is controlled to be between 0.5-1.0%.
  • Mn In the ultra-high-strength dual-phase steel of the present invention, Mn is an element that strongly improves the austenite hardenability, and it can effectively increase the strength of the steel by forming more martensite. Therefore, the mass percentage of Mn in the ultra-high-strength dual-phase steel of the present invention is controlled to be between 2.5-3.0%.
  • the mass percentage of Mn can be controlled between 2.5-2.8%.
  • Al is a deoxidizing element, which can deoxidize and refine grains in the steel. Therefore, the mass percentage of Al in the ultra-high-strength dual-phase steel of the present invention is controlled to be between 0.02-0.05%.
  • Nb and Ti are used as carbonitride precipitation elements, which can refine crystal grains and precipitate carbonitrides, improve the strength of the material, and can be added separately or in combination.
  • the mass percentage of Nb or Ti in the steel is higher than 0.05%, the strengthening effect is not significant. Therefore, in the ultra-high-strength dual-phase steel of the present invention, the mass percentage of Nb is controlled to be between 0.02-0.05%, and the mass percentage of Ti is controlled to be between 0.02-0.05%.
  • B In the ultra-high-strength dual-phase steel of the present invention, B is used as a strong hardenability element, and an appropriate amount of B can improve the hardenability of the steel and promote the formation of martensite. Therefore, the mass percentage of B in the ultra-high-strength dual-phase steel of the present invention is controlled to be between 0.001% and 0.003%.
  • the unavoidable impurities include P, S and N elements, and their content is controlled to at least one of the following items: P ⁇ 0.01%, S ⁇ 0.002%, N ⁇ 0.004%.
  • P, S and N elements are unavoidable impurity elements in the steel.
  • the mass percentage of each chemical element satisfies at least one of the following items:
  • the phase ratio (volume ratio) of the martensite is> 90%.
  • the martensite contains coherent distribution of epsilon carbides.
  • the ultra-high-strength dual-phase steel of the present invention its performance satisfies at least one of the following items: yield strength ⁇ 900 MPa, tensile strength ⁇ 1300 MPa, elongation after fracture ⁇ 5%, initial The hydrogen content is less than or equal to 10ppm; when the preset stress is greater than or equal to one time the tensile strength, the delayed cracking will not occur after being soaked in 1mol/L hydrochloric acid for 300 hours.
  • the ultra-high-strength dual-phase steel of the present invention its performance satisfies at least one of the following items: yield strength ⁇ 930MPa, tensile strength ⁇ 1320MPa, elongation after fracture ⁇ 5.5%, initial When the hydrogen content is less than or equal to 7ppm, and the pre-stress is 1.2 times the tensile strength, the delayed cracking will not occur after being soaked in 1mol/L hydrochloric acid for 300 hours.
  • the yield strength is greater than or equal to 930 MPa
  • the tensile strength is greater than or equal to 1320 MPa
  • the elongation after fracture is greater than or equal to 5.5%
  • the initial hydrogen content is less than or equal to 7 ppm.
  • another object of the present invention is to provide a method for manufacturing ultra-high-strength dual-phase steel, the yield strength of the ultra-high-strength dual-phase steel produced by the method is ⁇ 900MPa, the tensile strength ⁇ 1300MPa, and the elongation after fracture ⁇ 5%, initial hydrogen content ⁇ 10ppm; when the preset stress is greater than or equal to twice the tensile strength, no delayed cracking will occur after being soaked in 1mol/L hydrochloric acid for 300 hours. It can be effectively applied to automobile safety structural parts Manufacturing has good promotion and application value and prospects.
  • the present invention proposes the above-mentioned manufacturing method of ultra-high-strength dual-phase steel, which includes the following steps:
  • Annealing heating at a heating rate of 3-10°C/s to an annealing soaking temperature of 800-850°C, preferably 805-845°C, annealing time of 40-200s, and then a rapid rate of 30-80°C/s Cooling, the starting temperature of rapid cooling is 670 ⁇ 730°C;
  • tempering temperature is 260-320°C, preferably 260-310°C, and the tempering time is 100-400s, preferably 100-300s;
  • Annealing uses a combination of high temperature soaking + medium temperature tempering.
  • High-temperature soaking causes more austenite transformation to occur, and more martensite is obtained during the subsequent rapid cooling, which ultimately ensures higher strength before tempering.
  • medium-temperature tempering makes the yield ratio of the material moderate.
  • the yield ratio of the ultra-high-strength dual-phase steel of the present invention is between 0.70 and 0.75.
  • the relevant process parameters are controlled by using medium and low temperature tempering treatment.
  • the martensite can be easily precipitated during tempering.
  • Uniform, small, and dispersed coherent ⁇ carbides on the other hand, the method of long-term tempering at medium and low temperatures can remove the excess hydrogen in the steel plate to the greatest extent, so that it can diffuse out of the steel plate, so that the original state of the steel plate The hydrogen content is reduced.
  • it is beneficial to reduce the hardness of martensite and the diffusion of hydrogen in the steel plate it is also very beneficial to the mechanical properties and delayed cracking performance of the steel.
  • step (1) the continuous casting pulling speed is controlled to be 0.9-1.5 m/min during the continuous casting process.
  • step (1) continuous casting can adopt a large water volume secondary cooling mode for rapid cooling to minimize segregation.
  • step (2) the cast slab is controlled to be soaked at a temperature of 1220 to 1260°C, preferably 1220 to 1250°C; then rolling, and the final rolling temperature is controlled to be 880 to 1250°C. 920°C, cooling at a rate of 20-70°C/s after rolling; then coiling, the coiling temperature is 600-650°C, preferably 605-645°C, and heat preservation treatment is performed after coiling.
  • heat preservation treatment is performed after coiling, and heat preservation is performed for 1-5 hours.
  • the heating temperature in the step (2), in order to ensure the stability of the rolling load, the heating temperature is controlled to be above 1220°C, and at the same time, to prevent the increase of oxidation burning loss ,
  • the upper limit of the control heating temperature is 1260°C, therefore, the final control of the cast slab is soaked at a temperature of 1220-1260°C.
  • step (3) the cold rolling reduction ratio is controlled to be 45-65%.
  • the surface scale of the steel sheet can be removed by pickling.
  • step (6) the leveling reduction rate is controlled to be ⁇ 0.3%.
  • the leveling reduction rate is controlled to be ⁇ 0.3%.
  • the step (7) can be implemented by a conventional electro-galvanizing method.
  • double-sided plating is performed, and the weight of the plating layer on one side is in the range of 10-100 g/m 2.
  • the ultra-high-strength dual-phase steel and the manufacturing method thereof according to the present invention have the following advantages and beneficial effects:
  • the ultra-high-strength dual-phase steel of the present invention adopts reasonable composition design, adopts the design of medium Si and low Al, reduces the use of alloying elements such as Si and Al, and avoids the surface quality caused by high Si and the plate caused by high Al. Problems such as blank defects.
  • the steel does not contain precious alloy elements such as Cr and Mo, has a small alloy content, has good manufacturability, has good economic efficiency, and effectively controls alloy costs.
  • the yield strength of the ultra-high-strength dual-phase steel is ⁇ 900MPa, the tensile strength is ⁇ 1300MPa, the elongation after fracture is ⁇ 5%, and the initial hydrogen content is ⁇ 10ppm; when the preset stress is greater than or equal to one time the tensile strength, 1mol /L hydrochloric acid soaking for 300 hours does not cause delayed cracking, which can be applied to the manufacture of automobile safety structural parts, and has good promotion and application value and prospects.
  • the relevant process parameters are controlled by using medium and low temperature tempering treatment.
  • the martensite can be easily precipitated uniformly, finely, and Dispersed coherent ⁇ carbides
  • the method of long-term tempering at medium and low temperatures can remove the excess hydrogen in the steel plate to the greatest extent, and make it diffuse out of the steel plate, thereby reducing the hydrogen content of the original state of the steel plate.
  • it beneficial to reduce the hardness of martensite and the diffusion of hydrogen in the steel plate it is also very beneficial to the mechanical properties and delayed cracking performance of the steel. It effectively ensures that the ultra-high strength dual-phase steel produced has excellent mechanical properties and excellent delay resistance. The characteristics of cracking and low initial hydrogen content.
  • Figure 1 shows the structure of the cold rolled and annealed dual phase steel of Example 1.
  • Table 1 lists the mass percentages of various chemical elements in the steel grades corresponding to the ultra-high-strength dual-phase steels of Examples 1-7 and the steels of Comparative Examples 1-14.
  • the ultra-high-strength dual-phase steels of Examples 1-7 and the steels of Comparative Examples 1-14 of the present invention were prepared by the following steps:
  • Hot rolling control the cast slab to soak at a temperature of 1220 ⁇ 1260°C; then roll, control the final rolling temperature to be 880 ⁇ 920°C, cool it at a rate of 20 ⁇ 70°C/s after rolling; then carry out coiling ,
  • the coiling temperature is 600 ⁇ 650°C, and the heat preservation cover is used for heat preservation after coiling;
  • Annealing heating at a heating rate of 3-10°C/s to an annealing soaking temperature of 800-850°C, annealing time of 40-200s, and then rapid cooling at a rate of 30-80°C/s, the beginning of rapid cooling
  • the temperature is 670 ⁇ 730°C;
  • Tempering temperature is 260 ⁇ 320°C, and tempering time is 100 ⁇ 400s;
  • Double-sided electro-galvanization the weight of the single-sided coating is 10-100g/m 2 .
  • the chemical composition and related process parameters of the ultra-high-strength dual-phase steel of Examples 1-7 all meet the control requirements of the design specification of the present invention.
  • the chemical compositions of the steels of Comparative Examples 1-6 all have parameters that fail to meet the design requirements of the present invention; although the chemical composition of the N steel grades corresponding to Comparative Examples 7-14 meets the design requirements of the present invention, the relevant process parameters do not exist. Meet the parameters of the design specification of the present invention.
  • Table 2-1 and Table 2-2 list the specific process parameters of the ultra-high-strength dual-phase steel of Example 1-7 and the steel of Comparative Example 1-14.
  • the performance test method refers to the GB/T 13239-2006 low-temperature tensile test method for metallic materials, prepares standard specimens, performs static stretching on a tensile testing machine, and obtains the corresponding stress-strain curve. After data processing, the yield strength and resistance are finally obtained. Tensile strength and elongation at break parameters.
  • the measurement method of hydrogen content Heat the sample to a certain temperature, and use a hydrogen analyzer to measure the concentration of hydrogen released with the temperature change (increase) to determine the initial hydrogen content in the steel.
  • Table 3 lists the performance test results of the ultra-high-strength dual-phase steels of Examples 1-7 and the steels of Comparative Examples 1-14.
  • high-strength steel with a strength of 1300Mpa or more can be produced.
  • the yield strength of each embodiment of the present invention is ⁇ 900MPa
  • the tensile strength is ⁇ 1300MPa
  • the elongation after fracture is ⁇ 5%.
  • the initial hydrogen content is ⁇ 10ppm.
  • the super-strength dual-phase steel of each embodiment has super-high strength and delayed cracking performance that is significantly better than that of the comparable steel grades of the same level. When the pre-stress Soaked in hydrochloric acid for 300 hours without delayed cracking.
  • the ultra-high-strength dual-phase steel of each embodiment has excellent performance, can be applied to the manufacture of automobile safety structural parts, and has good promotion and application value and prospects.

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Abstract

本发明公开了一种超高强双相钢，其基体组织为铁素体+马氏体，其中铁素体和马氏体呈岛状均匀分布，超高强双相钢含有质量百分比如下的下述化学元素：C：0.12-0.2％，Si：0.5-1.0％，Mn：2.5-3.0％，Al：0.02-0.05％，Nb：0.02-0.05％，Ti：0.02-0.05％，B：0.001％-0.003％。本发明还公开了上述超高强双相钢的制造方法，包括步骤：冶炼和连铸、热轧、冷轧、退火、回火和平整。本发明超高强双相钢不仅有较好的力学性能，还具有优异的耐延迟开裂性和较低初始氢含量，可以适用于汽车安全结构件的制造。

Description

一种超高强双相钢及其制造方法 技术领域
本发明涉及一种金属材料及其制造方法，尤其涉及一种双相钢及其制造方法。
背景技术
随着汽车工业轻量化减重和安全性的需要，市场对强度更高的钢板需求量越来越多。其中,双相钢由于生产成本低、可制造性强同时具有低屈服强度、高抗拉强度以及高的初始加工硬化速率等优良的性能在汽车零部件中得到了广泛地使用。
目前市场上强度等级需求主要以80公斤、100公斤级的为主，目前最高强度级别为1180DP牌号，其抗拉强度大于等于1200MPa，屈服强度约为850MPa，总延伸率约为10％。冷轧双相钢的生产采用临界区连续退火工艺，其抗拉强度由退火组织中的马氏体分数所决定，马氏体分数越高则抗拉强度越高，这就要求在生产时需要采用较高的退火温度以形成更多的马氏体分数，目前所能商业化生产的双相钢最高强度等级为1180MPa，即DP 1180钢。
公开号为CN109504930A，公开日为2019年3月22日，名称为“抗拉强度大于1300MPa的热镀锌钢板及其生产方法”的中国专利文献，公开了一种抗拉强度大于1300MPa的热镀锌钢板及其生产方法，所述热镀锌钢板基板化学成分组成及其质量百分含量为：C：0.1～0.2％，Mn：1.3～2.0％，S≤0.005％，P≤0.02％，Si：0.2～0.3％，Als：0.4～1.0％，Nb：0.01～0.03％，Ti：0.04～0.08％，B：0.001～0.004％，Mo：0.2～0.3％，Cr：0.05～0.10％，V：0.01～0.02％，余量为Fe以及不可避免的杂质。所述板坯加热工序，加热温度1200～1320℃，加热时间120～200min；所述热轧工序，粗轧轧制3～7道次；精轧进口温度1020～1080℃，终轧温度820～880℃；卷取温度550～650℃；所述生产方法包括板坯加热、热轧、酸轧、连续热镀锌、光整和钝化工序；所述连续热镀锌工序，均热温度为760～840℃、保温时间100～200s，缓冷温度680～740℃、缓冷冷却速率10～20℃/s，快冷温度420～450℃、快冷冷却速率35～65℃/s，镀锌温度458～462℃、镀锌时间5～15s。
公开号为CN108486494A，公开日为2018年9月4日，名称为“钒微合金化 1300MPa级别高强热轧钢板和冷轧双相钢板的生产方法”的中国专利文献，公开了一种钒微合金化1300MPa级别高强热轧钢板和冷轧双相钢板的生产方法，其化学成分为：0.10-0.30wt％C，1.50-4.50wt％Mn，0.00-0.120wt％Al，0.00-0.90wt％Si，0.05-0.50％V，P≤0.020wt％，S≤0.02wt％，Fe:余量。该高强钢通将纳米碳化钒粒子析出强化与马氏体相变强化相结合，显著提高了现有双相钢的强度，同时还保证了较高的生产效率。
公开号为CN109628846A，公开日为2019年4月16日，名称为“1300MPa级汽车用超高强度冷轧钢板及其生产方法”的中国专利文献，公开了一种热成型钢板及制造方法，其化学成分为：C：0.1～0.2％，Mn：1.3～2.0％，S≤0.005％，P≤0.02％，Si：0.2～0.3％，Als：0.4～1.0％，Nb：0.01～0.03％，Ti：0.04～0.08％，B：0.001～0.004％，Mo：0.2～0.3％，Cr：0.05～0.10％，V：0.01～0.02％，Fe:余量。所述生产方法包括炼钢、连铸、热轧、酸轧、连续退火、平整拉矫工序；所述热轧工序，板坯加热温度≥1200℃，粗轧轧制3～7道次、粗轧后中间坯厚度28～40mm，精轧进口温度1020～1100℃，终轧温度820～900℃，卷取温度550～650℃；所述酸轧工序，酸洗后进行冷轧，冷轧压下率≥45％，所述连续退火工序，均热段保温温度为760～840℃，保温时间为60～225s；过时效段保温温度为250～320℃，过时效段保温时间为300～1225s。
由此可见，现有专利文献涉及到产品抗拉强度等级大于或等于1300MPa的主要产品以镀锌为主，且部分专利含有高Si、高Al，不利于表面质量以及生产制造。在部分专利技术中含有较多Cr、Mo等贵合金元素，生产成本较高。
发明内容
本发明的目的之一在于提供一种超高强双相钢，该超高强双相钢通过合理的化学元素成分设计，采用中Si低Al的设计，减少Si、Al等合金元素的使用，避免因高Si带来的表面质量及高Al带来的板坯缺陷等问题。
此外本发明超高强双相钢中不采用Cr、Mo等贵合金元素，有效控制了合金成本，同时降低杂质元素P含量、S含量，有利于性能提升、延迟开裂的改善。该超高强双相钢的屈服强度≥900MPa、优选≥930MPa，抗拉强度≥1300MPa、优选大≥1320MPa，断后伸长率≥5％、优选≥5.5％，起始氢含量≤10ppm、优选≤7ppm；在预置应力大于等于一倍抗拉强度的情况下，以1mol/L的盐酸浸泡300小时不发生延 迟开裂，优选在预置应力为抗拉强度1.2倍的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂，可以有效适用于汽车安全结构件的制造，具有良好的推广应用价值和前景。
为了实现上述目的，本发明提供了一种超高强双相钢，其基体组织为铁素体+马氏体，其中铁素体和马氏体呈岛状均匀分布，所述超高强双相钢除了Fe以外还含有质量百分比如下的下述化学元素：
C：0.12-0.2％，Si：0.5-1.0％，Mn：2.5-3.0％，Al：0.02-0.05％，Nb：0.02-0.05％，Ti：0.02-0.05％，B：0.001％-0.003％。
进一步地，在本发明所述的超高强双相钢中，其各化学元素质量百分比为：
C：0.12-0.2％，Si：0.5-1.0％，Mn：2.5-3.0％，Al：0.02-0.05％，Nb：0.02-0.05％，Ti：0.02-0.05％，B：0.001％-0.003％，余量为Fe和其他不可避免的杂质。
在本发明所述的超高强双相钢中，各化学元素的设计原理如下所述：
C：在本发明所述的超高强双相钢中，C是固溶强化元素，是材料获得高强度的保证。但是，需要注意的是，钢中含C量越高，马氏体越硬，发生延迟开裂的倾向越大。因此产品设计时，尽量选择低碳的设计，在本发明所述的超高强双相钢中控制C的质量百分比在0.12-0.2％之间。
在一些优选的实施方式中，C的质量百分比可以控制在0.14-0.18％之间。
Si：在本发明所述的超高强双相钢中，Si在钢中起到提高延伸率的作用。Si对钢的组织影响也很大，促进铁素体的纯净化和残余奥氏体的形成。同时能提高马氏体的抗回火性能，可以抑制Fe 3C的析出和长大，从而使回火时，形成的析出物以ε碳化物为主。但需要注意的是，当钢中Si的质量百分比低于0.5％，会影响钢的延伸率及抗回火性能，而若Si的质量百分比高于1.0％，则会带来其它的冶金质量缺陷。因此，在本发明所述的超高强双相钢中控制Si的质量百分比在0.5-1.0％之间。
Mn：在本发明所述的超高强双相钢中，Mn是强烈提高奥氏体淬透性的元素，其可以通过形成更多的马氏体从而有效提高钢的强度。因此，在本发明所述的超高强双相钢中控制Mn的质量百分比在2.5-3.0％之间。
在一些优选的实施方式中，Mn的质量百分比可以控制在2.5-2.8％之间。
Al：在本发明所述的超高强双相钢中，Al是脱氧元素，其可以在钢中脱氧作用和细化晶粒的作用。因此，在本发明所述的超高强双相钢中控制Al的质量百分比在0.02-0.05％之间。
Nb和Ti：在本发明所述的超高强双相钢中，Nb和Ti作为碳氮化物析出元素，可以细化晶粒和析出碳氮化物，提高材料的强度，可以单独添加或复合添加。但是，需要注意的是，若钢中Nb或Ti的质量百分含量高于0.05％，强化作用不显著。因此，在本发明所述的超高强双相钢中控制Nb的质量百分比在0.02-0.05％之间，控制Ti的质量百分比在0.02-0.05％之间。
B：在本发明所述的超高强双相钢中，B作为强淬透性元素，适量的B可以提高钢的淬透性，促进马氏体的形成。因此在本发明所述的超高强双相钢中控制B的质量百分比在0.001％-0.003％之间。
进一步地，在本发明所述的超高强双相钢中，其中不可避免的杂质包括P、S和N元素，其含量控制为下述各项的至少其中之一：P≤0.01％，S≤0.002％，N≤0.004％。
上述技术方案中，在本发明所述的超高强双相钢中，P、S和N元素均是钢中不可避免的杂质元素，在钢中P、S和N元素含量越低越好。S易形成MnS夹杂物，严重影响扩孔率；P元素会降低钢的韧性，对延迟开裂不利；钢中N元素含量过高，容易导致板坯表面裂纹，大大影响钢的性能。因此，在本发明所述的超高强双相钢中，控制P的质量百分比为P≤0.01％，控制S的质量百分比为S≤0.002％，控制N的质量百分比为N≤0.004％。
进一步地，在本发明所述的超高强双相钢中，其各化学元素质量百分含量满足下述各项的至少其中之一：
C：0.14-0.18％，
Mn：2.5-2.8％。
进一步地，在本发明所述的超高强双相钢中，所述马氏体的相比例(体积比)＞90％。
进一步地，在本发明所述的超高强双相钢中，所述马氏体中含有共格分布的ε碳化物。
进一步地，在本发明所述的超高强双相钢中，其性能满足下述各项的至少其中之一：屈服强度≥900MPa，抗拉强度≥1300MPa，断后伸长率≥5％，起始氢含量≤10ppm；在预置应力大于等于一倍抗拉强度的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂。
进一步地，在本发明所述的超高强双相钢中，其性能满足下述各项的至少其中之一：屈服强度≥930MPa，抗拉强度≥1320MPa，断后伸长率≥5.5％，起始氢含量 ≤7ppm，且在预置应力为抗拉强度1.2倍的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂。
优选地，在本发明所述的超高强双相钢中，屈服强度≥930MPa，抗拉强度≥1320MPa，断后伸长率≥5.5％，起始氢含量≤7ppm，且在预置应力为抗拉强度1.2倍的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂。
相应地，本发明的另一目的在于提供一种超高强双相钢的制造方法，采用该制造方法制得的超高强双相钢的屈服强度≥900MPa，抗拉强度≥1300MPa，断后伸长率≥5％，起始氢含量≤10ppm；在预置应力大于等于一倍抗拉强度的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂，其可以有效适用于汽车安全结构件的制造，具有良好的推广应用价值和前景。
为了实现上述目的，本发明提出了上述的超高强双相钢的制造方法，包括步骤：
(1)冶炼和连铸；
(2)热轧；
(3)冷轧；
(4)退火：以3-10℃/s的加热速度升温到退火均热温度800～850℃、优选805～845℃，退火时间为40～200s，然后以30～80℃/s的速度快速冷却，快速冷却的开始温度为670～730℃；
(5)回火：回火温度为260～320℃、优选260～310℃，回火时间为100～400s、优选100～300s；
(6)平整；
(7)电镀锌。
退火采用的为高温均热+中温回火相结合的方式。高温均热使得发生较多的奥氏体转变，在随后的快速冷却时得到更多的马氏体，最终保证回火前有更高的强度，中温回火一方面使得材料的屈强比适中，另一方面也可以起到更好的改善延迟开裂的效果。在优选的实施方案中，本发明所述的超高强双相钢的屈强比在0.70-0.75之间。
在本发明所述的超高强双相钢的制造方法中，在连续退火后，通过采用中低温回火处理，对相关工艺参数进行控制，一方面可以使马氏体在回火时，易于析出均匀、细小、弥散的共格型ε碳化物，另一方面中低温长时回火的方式，可以最大程度的去除钢板中过剩的氢，使之扩散到钢板外，从而使钢板的原始状态的氢含量降 低。不仅有利于降低马氏体的硬度及钢板内部氢的扩散，还对钢的力学性能及延迟开裂性能十分有利。
进一步地，在本发明所述的制造方法中，在步骤(1)中，连铸过程中控制连铸拉速为0.9-1.5m/min。
在上述技术方案中，在本发明所述的制造方法中，在步骤(1)中连铸可以采用大水量二冷模式进行快速冷却，尽量减小偏析。
进一步地，在本发明所述的制造方法中，在步骤(2)中，控制铸坯以1220～1260℃、优选1220～1250℃的温度均热；然后轧制，控制终轧温度为880～920℃，轧后以20～70℃/s的速度冷却；然后进行卷取，卷取温度为600～650℃、优选605～645℃，卷取后进行保温处理。优选地，卷取后进行保温处理，保温1-5小时。
在本发明所述的超高强双相钢的制造方法中，在所述步骤(2)中，为保证轧制负荷的稳定，控制加热温度在1220℃以上，同时为防止氧化烧损的增大，控制加热温度的上限为1260℃，因此，最终控制铸坯以1220～1260℃的温度均热。
进一步地，在本发明所述的制造方法中，在步骤(3)中，控制冷轧压下率为45～65％。
上述方案中，在所述步骤(3)中，在控制冷轧压下率为45～65％冷轧前，可以通过酸洗去除钢板表面氧化铁皮。
进一步地，在本发明所述的制造方法中，在步骤(6)中，控制平整压下率≤0.3％。
在本发明上述方案中，在所述步骤(6)中，为保证钢板的平整度，需要进行一定的平整量，然而过大的平整量会使得钢的屈服强度上升较多。因此，在本发明所述的制造方法中，控制平整压下率≤0.3％。
在本发明上述方案中，可采用常规的电镀锌的方法实施所述步骤(7)。优选地，进行双面镀，且单面的镀层重量在10-100g/m 2的范围内。
本发明所述的超高强双相钢及其制造方法相较于现有技术具有如下所述的优点以及有益效果：
本发明所述的超高强双相钢采用合理的成分设计，采用中Si低Al的设计，减少Si、Al等合金元素的使用，避免因高Si带来的表面质量及高Al带来的板坯缺陷等问题。此外，钢中不含Cr、Mo等贵合金元素，合金含量少，可制造性好，具有良好的经济性，有效控制了合金成本。该超高强双相钢的屈服强度≥900MPa，抗拉强度≥1300MPa，断后伸长率≥5％，起始氢含量≤10ppm；在预置应力大于等于一倍 抗拉强度的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂，可以适用于汽车安全结构件的制造，具有良好的推广应用价值和前景。
此外，在本发明所述的制造方法中，在连续退火后，通过采用中低温回火处理，对相关工艺参数进行控制，一方面可以使马氏体在回火时，易于析出均匀、细小、弥散的共格型ε碳化物，另一方面中低温长时回火的方式，可以最大程度的去除钢板中过剩的氢，使之扩散到钢板外，从而使钢板的原始状态的氢含量降低。不仅有利于降低马氏体的硬度及钢板内部氢的扩散，还对钢的力学性能及延迟开裂性能十分有利，有效保证了制得的超高强双相钢具有优异的力学性能、优异的耐延迟开裂性和较低的初始氢含量的特性。
附图说明
图1显示实施例1冷轧退火双相钢的组织。
具体实施方式
下面将结合具体的实施例对本发明所述的超高强双相钢及其制造方法做进一步的解释和说明，然而该解释和说明并不对本发明的技术方案构成不当限定。
实施例1-7和对比例1-14
表1列出了实施例1-7的超高强双相钢和对比例1-14钢对应的钢种中各化学元素质量百分比。
表1(wt％，余量为Fe和其他除了P、S以及N以外的不可避免的杂质)
Figure PCTCN2021095807-appb-000001
Figure PCTCN2021095807-appb-000002
本发明所述实施例1-7的超高强双相钢和对比例1-14的钢均采用以下步骤制得：
(1)冶炼和连铸：其中在连铸过程中，控制连铸拉速为0.9-1.5m/min，并采用大水量二冷模式进行快速冷却；
(2)热轧：控制铸坯以1220～1260℃的温度均热；然后轧制，控制终轧温度为880～920℃，轧后以20～70℃/s的速度冷却；然后进行卷取，卷取温度为600～650℃，卷取后采用保温罩进行保温处理；
(3)冷轧：控制冷轧压下率为45～65％；
(4)退火：以3-10℃/s的加热速度升温到退火均热温度800～850℃，退火时间为40～200s，然后以30～80℃/s的速度快速冷却，快速冷却的开始温度为670～730℃；
(5)回火：回火温度为260～320℃，回火时间为100～400s；
(6)平整：控制平整压下率≤0.3％；
(7)双面电镀锌，单面镀层重量10-100g/m 2。
需要说明的是，实施例1-7的超高强双相钢的化学成分和相关工艺参数均满足本发明设计规范控制要求。对比例1-6的钢化学成分均存在未能满足本发明设计的要求的参数；对比例7-14对应的N钢种的化学成分虽然满足本发明设计要求，但是相关工艺参数均存在未能满足本发明设计规范的参数。
表2-1和表2-2列出了实施例1-7的超高强双相钢和对比例1-14钢的具体工艺参数。
表2-1
Figure PCTCN2021095807-appb-000003
Figure PCTCN2021095807-appb-000004
表2-2
Figure PCTCN2021095807-appb-000005
Figure PCTCN2021095807-appb-000006
将实施例1-7的超高强双相钢和对比例1-14钢进行各项性能测试，所得的测试结果列于表3中。
性能测试方法参照GB/T 13239-2006金属材料低温拉伸试验方法，制备标准试样，在拉伸试验机上进行静态拉伸，得到相应应力-应变曲线，经过数据处理，最终得到屈服强度、抗拉强度和断裂延伸率参数。
氢含量的测量方法：将样品加热至一定温度，利用氢分析仪测量通过随温度变化(升高)释放出的氢的浓度，从而判断钢中的起始氢含量。
表3列出了实施例1-7的超高强双相钢和对比例1-14钢的性能测试结果。
表3
Figure PCTCN2021095807-appb-000007
Figure PCTCN2021095807-appb-000008
注：钢板在一定内应力水平下浸泡在1mol/L的盐酸中300小时的结果：Ο表示未开裂，X表示开裂。
由表3可看出，按照本发明可以制造出强度1300Mpa以上的高强度钢，本发明各实施例的屈服强度均≥900MPa，抗拉强度均≥1300MPa，断后伸长率均≥5％，起始氢含量均≤10ppm。各实施例的超强双相钢均具有超高的强度和明显优于同等级别的对比钢种的延迟开裂性能，其在预置应力大于等于一倍抗拉强度的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂。各实施例的超高强双相钢的性能十分优异，可以适用于汽车安全结构件的制造，具有良好的推广应用价值和前景。
需要说明的是，本发明的保护范围中现有技术部分并不局限于本申请文件所给出的实施例，所有不与本发明的方案相矛盾的现有技术，包括但不局限于在先专利文献、在先公开出版物，在先公开使用等等，都可纳入本发明的保护范围。此外，本案中各技术特征的组合方式并不限本案权利要求中所记载的组合方式或是具体实施例所记载的组合方式，本案记载的所有技术特征可以以任何方式进行自由组合或结合，除非相互之间产生矛盾。
还需要注意的是，以上所列举的实施例仅为本发明的具体实施例。显然本发明不局限于以上实施例，随之做出的类似变化或变形是本领域技术人员能从本发明公开的内容直接得出或者很容易便联想到的，均应属于本发明的保护范围。

Claims (15)

  1. 一种超高强双相钢，其特征在于，其基体组织为铁素体+马氏体，其中铁素体和马氏体呈岛状均匀分布，所述超高强双相钢除了Fe以外还含有质量百分比如下的下述化学元素：
    C：0.12-0.2％，Si：0.5-1.0％，Mn：2.5-3.0％，Al：0.02-0.05％，Nb：0.02-0.05％，Ti：0.02-0.05％，B：0.001％-0.003％。
  2. 如权利要求1所述的超高强双相钢，其特征在于，其各化学元素质量百分比为：
    C：0.12-0.2％，Si：0.5-1.0％，Mn：2.5-3.0％，Al：0.02-0.05％，Nb：0.02-0.05％，Ti：0.02-0.05％，B：0.001％-0.003％，余量为Fe和其他不可避免的杂质。
  3. 如权利要求2所述的超高强双相钢，其特征在于，其中不可避免的杂质包括P、S和N元素，其含量控制为下述各项的至少其中之一：P≤0.01％，S≤0.002％，N≤0.004％。
  4. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，其各化学元素满足下述各项的至少其中之一：
    C：0.14-0.18％，
    Mn：2.5-2.8％。
  5. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，所述马氏体的相比例＞90％。
  6. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，所述马氏体中含有共格分布的ε碳化物。
  7. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，其性能满足下述各项的至少其中之一：屈服强度≥900MPa，抗拉强度≥1300MPa，断后伸长率≥5％，起始氢含量≤10ppm；在预置应力大于等于一倍抗拉强度的情况下，在1mol/L的盐酸浸泡300小时不发生延迟开裂。
  8. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，其性能满足下述各项：屈服强度≥930MPa，抗拉强度≥1320MPa，断后伸长率≥5.5％，起始氢含量≤7ppm，且在预置应力为抗拉强度1.2倍的情况下，以1mol/L的盐酸浸泡300小时不发生延迟开裂。
  9. 如权利要求1-3中任一项所述的超高强双相钢，其特征在于，所述超高强双相 钢的屈强比在0.70-0.75之间。
  10. 一种如权利要求1-9中任意一项所述的超高强双相钢的制造方法，其特征在于，包括步骤：
    (1)冶炼和连铸；
    (2)热轧；
    (3)冷轧；
    (4)退火：以3-10℃/s的加热速度升温到退火均热温度800～850℃，退火时间为40～200s，然后以30～80℃/s的速度快速冷却，快速冷却的开始温度为670～730℃；
    (5)回火：回火温度为260～320℃，回火时间为100～400s；
    (6)平整；
    (7)电镀。
  11. 如权利要求10所述的制造方法，其特征在于，在步骤(1)中，连铸过程中控制连铸拉速为0.9-1.5m/min。
  12. 如权利要求10所述的制造方法，其特征在于，在步骤(2)中，控制铸坯以1220～1260℃的温度均热；然后轧制，控制终轧温度为880～920℃，轧后以20～70℃/s的速度冷却；然后进行卷取，卷取温度为600～650℃，卷取后进行保温处理。
  13. 如权利要求10所述的制造方法，其特征在于，在步骤(3)中，控制冷轧压下率为45～65％。
  14. 如权利要求10所述的制造方法，其特征在于，在步骤(6)中，控制平整压下率≤0.3％；和/或，步骤(7)中，双面电镀锌，单面镀层重量10-100g/m 2。
  15. 利要求10所述的制造方法，其特征在于，步骤(2)中，控制铸坯以1220～1250℃的温度均热，卷取温度为605～645℃；步骤(4)中，退火均热温度为805～845℃；步骤(5)中，回火温度为260～310℃，回火时间为优选100～300s。
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