WO2019001423A1 - 一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法 - Google Patents

一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法 Download PDF

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WO2019001423A1
WO2019001423A1 PCT/CN2018/092876 CN2018092876W WO2019001423A1 WO 2019001423 A1 WO2019001423 A1 WO 2019001423A1 CN 2018092876 W CN2018092876 W CN 2018092876W WO 2019001423 A1 WO2019001423 A1 WO 2019001423A1
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formability
cost
steel sheet
low
mpa
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PCT/CN2018/092876
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English (en)
French (fr)
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薛鹏
朱晓东
李伟
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宝山钢铁股份有限公司
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Priority to KR1020197038709A priority Critical patent/KR20200012953A/ko
Priority to EP18825106.0A priority patent/EP3647454A4/en
Priority to JP2019568097A priority patent/JP6924284B2/ja
Priority to US16/624,270 priority patent/US11339451B2/en
Publication of WO2019001423A1 publication Critical patent/WO2019001423A1/zh

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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
    • 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
    • C21D8/0236—Cold 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
    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21C—MANUFACTURE OF METAL SHEETS, WIRE, RODS, TUBES OR PROFILES, OTHERWISE THAN BY ROLLING; AUXILIARY OPERATIONS USED IN CONNECTION WITH METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL
    • B21C37/00—Manufacture of metal sheets, bars, wire, tubes or like semi-manufactured products, not otherwise provided for; Manufacture of tubes of special shape
    • B21C37/02—Manufacture of metal sheets, bars, wire, tubes or like semi-manufactured products, not otherwise provided for; Manufacture of tubes of special shape of sheets
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/005—Heat treatment of ferrous alloys containing Mn
    • 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
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/02—Hardening by precipitation
    • 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
    • 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
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
    • 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
    • C21D8/021—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular fabrication or treatment of ingot or slab
    • C21D8/0215—Rapid solidification; Thin strip casting
    • 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
    • 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
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • 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
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/0062—Heat-treating apparatus with a cooling or quenching zone
    • 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
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • 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
    • 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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • 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
    • C21D2211/005—Ferrite
    • 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 relates to a dual-phase steel plate and a manufacturing method thereof, in particular to a low-cost and high-formability 1180 MPa grade cold-rolled annealing double-phase steel plate and a manufacturing method thereof, the steel plate having a yield strength of 850 MPa or more and a tensile strength of 1180 MPa or more. Molding properties, especially excellent cold bending properties.
  • High-strength dual-phase steel has good mechanical properties and performance, and is suitable for the manufacture of structural parts.
  • a conventional cold-rolled ultra-high-strength dual-phase steel such as the Chinese Patent Application Publication No. CN102227511B, discloses a high-strength cold-rolled steel sheet having a TS of 1180 MPa or more and excellent formability such as hole expandability and flexibility, and high-strength hot-dip galvanizing.
  • Steel sheets and methods for their production C: 0.05-0.3%, Si: 0.5-2.5%, Mn: 1.5-3.5%, P: 0.001-0.05%, S: 0.0001-0.01%, Al: 0.001-0.1%, N: 0.0005-0.01%, Cr 1.5% or less and including 0%.
  • the area ratio of the martensite phase in the entire structure is 30% or more (area occupied by the martensite phase) / (the proportion of the ferrite phase)
  • the area is greater than 0.45 and less than 1.5. Its uniqueness is that it relies on a high proportion of ferrite content to obtain better formability.
  • CiaS yield strength
  • TS tensile strength
  • Strength steel plate and its manufacturing method C: 0.05-0.15%, Si: 0.01-1.00%, Mn: 1.5-4.0%, P: 0.1% or less, S: 0.02% or less, Al: 0.01-0.5%, Nb: 0.005-0.1%, Ti: 0.005-0.1%, Cr: 0.01-2.0%, B: 0.0005-0.005%. It is required to set the average crystal grain size of ferrite and martensite to 10 ⁇ m or less. Its uniqueness lies in the refinement of crystal grains and the improvement of hardenability by the addition of B element and the corresponding process, and the addition of Cr improves the hardenability.
  • the steel sheet disclosed in Chinese Patent Application Publication No. CN102828106A has a tensile strength of 1180 MPa or more and is excellent in workability and low-temperature brittleness.
  • the high-strength steel sheet of the present invention contains C: 0.10 to 0.30%, Si: 1.40 to 3.0%, Mn: 0.5 to 3.0%, P: 0.1% or less, S: 0.05% or less, and Al: 0.005 to 0.20%, and N: 0.01% or less, O: 0.01% or less, the balance is composed of Fe and unavoidable impurities, and the volume ratio of the mixed structure of the primary martensite and the retained austenite (MA structure) is 6% with respect to the entire structure.
  • the volume fraction of retained austenite with respect to all the structures is 5% or more. It is unique in that it uses Si: 1.40 to 3.0% addition to obtain retained austenite in the final structure and improve molding properties.
  • Chinese Patent Application Publication No. CN106661701A is a method for producing a high-strength steel sheet by heat-treating a steel sheet having a yield strength YS>850 MPa, a tensile strength TS>1180 MPa, a total elongation of >13%, and an expansion.
  • the final tissue can contain between 3% and 15% retained austenite. Its uniqueness is that the addition of 1.2% ⁇ Si ⁇ 1.8%, the retained austenite is obtained in the final structure, and the molding property is improved.
  • Chinese Patent Application Publication No. CN106661653A is a method for producing a steel sheet having a yield strength YS greater than 1000 MPa, a tensile strength TS greater than 1150 MPa, and a total elongation E greater than 8%, the method comprising the steps of: preparing a steel sheet by rolling steel to The steel comprises 0.19% to 0.22% C, 2% to 2.6% Mn, 1.45% to 1.55% Si, 0.15% to 0.4% Cr, less than 0.020% P, less than 0.05% by weight.
  • the steel structure contains more than 80% tempered martensite and more than 5% retained austenite.
  • the object of the present invention is to provide a low-cost and high-formability 1180 MPa grade cold-rolled annealing double-phase steel plate and a manufacturing method thereof, and ensure the steel plate to reach a strength of 1180 MPa at a low cost by rational design of the alloy elements and the manufacturing process; Uniform martensite + ferrite dual phase structure to ensure excellent elongation and cold bending performance, and good formability; the yield strength of the duplex steel plate is greater than 850 MPa, and its tensile strength is greater than 1180 MPa, and its extension The rate is ⁇ 8%, and the 90 degree cold bending performance characterizes the parameter R/t ⁇ 2.5, where R represents the bending radius and t represents the plate thickness in mm.
  • the steel component of the invention is designed with a C+Mn-based component system to ensure a strength of 1180 MPa, and the steel does not add precious alloy elements such as Mo and Cr to ensure low cost.
  • the micro-addition of Nb and Ti achieves the effect of suppressing austenite grain growth and effectively refines crystal grains.
  • the special component design without adding Mo and Cr is combined with a special hot rolling low temperature coiling process to ensure that the hot coil tensile strength is less than 1000 MPa, and the manufacturability of cold rolling is ensured.
  • a special hot rolling coiling temperature is adopted: a low temperature coiling (400-500 ° C) of the bainite phase transformation zone in the hot rolling process, Ensure that the hot rolled microstructure has a bainite content of 80% or more; after the coiling, the cooling rate is not more than 0.3 °C/s to 370 °C, and does not enter the martensite transformation zone, ensuring that the hot rolling strength meets the requirements of cold rolling manufacturability.
  • the cold-rolled structure is a uniform deformed bainite structure with defects such as a large number of dislocations, and provides a large number of austenite nucleation points, so that the austenite exhibits explosive nucleation, and therefore austenite
  • the grains are significantly refined.
  • these defects have become a channel for high-speed diffusion of carbon atoms, and because the entire matrix structure is uniform, the difference in carbon concentration is small, so that austenite is rapidly formed, and the volume fraction is increased to obtain a fine uniform final structure.
  • the low-cost and high-formability 1180 MPa grade cold-rolled annealed duplex steel sheet of the present invention has a chemical composition mass percentage of C: 0.1% to 0.125%, Si: 0.4% to 0.8%, and Mn: 2.6%. 2.9%, Al: 0.01% to 0.05%, Nb: 0.01 to 0.03%, Ti: 0.01 to 0.03%, and the balance being Fe and inevitable impurities.
  • the microstructure of the dual-phase steel sheet of the present invention is a fine and uniform martensite + ferrite two-phase structure, wherein at least 75% of martensite is contained, and the rest is ferrite; the grain diameter of martensite
  • the grain size of the ferrite is not more than 5 ⁇ m and is not more than 5 ⁇ m.
  • the yield strength of the dual-phase steel plate of the invention is greater than 850 MPa; the tensile strength is greater than 1180 MPa; the elongation is ⁇ 8%; the 90-degree cold bending performance character is R/t ⁇ 2.5, wherein R represents the bending radius and t represents the thickness of the plate. , unit mm.
  • the addition of the C element serves to increase the strength of the steel and increase the hardness of the martensite. Therefore, the mass percentage of C is selected to be between 0.1% and 0.125% because: when the mass percentage of C is less than 0.1%, the strength of the steel sheet is affected, and it is disadvantageous to the formation amount and stability of austenite; When the mass percentage of C is higher than 0.125%, the martensite hardness is too high, and the grain size is coarse, which is disadvantageous for the forming property of the steel sheet. It is preferably 0.11% to 0.125%.
  • Si Adding Si can improve hardenability. Moreover, the solid solution of Si in steel can affect the interaction of dislocations and increase the work hardening rate. The elongation can be appropriately increased in the duplex steel, which is beneficial to obtain better formability.
  • the Si content is controlled to be Si: 0.4% to 0.8%, preferably 0.5% to 0.7%.
  • Mn The addition of Mn element is beneficial to the improvement of the hardenability of the steel and the strength of the steel sheet.
  • the mass percentage of Mn is 2.6% to 2.9% because the strength of the steel sheet is insufficient when the mass percentage of Mn is less than 2.6%; when the mass percentage of Mn is higher than 2.9%, the strength of the steel sheet is too high.
  • the molding properties are degraded and segregation is liable to occur. Therefore, the mass percentage of Mn controlled in the low-cost, high-formability 1180 MPa grade cold-rolled annealed duplex steel sheet according to the present invention is Mn: 2.6 to 2.9%, preferably 2.7% to 2.85%.
  • Al The addition of Al serves to deoxidize and refine the crystal grains. Therefore, the mass percentage of Al is controlled to be 0.01% to 0.05%, preferably 0.015 to 0.045%.
  • Nb 0.01 to 0.03% of Nb is added because: after adding a small amount of strong carbide forming element Nb to the microalloyed steel, the strain-induced precipitation phase passes through the pinning of the particle and the subgrain boundary during the controlled rolling process. However, the recrystallization temperature of the deformed austenite is reduced significantly, and the nucleation point is provided, which has obvious effect on the refinement of crystal grains; in the continuous austenitization process, the soaking carbon and nitride material points will pass through The particle pinning grain boundary mechanism prevents coarsening of the soaked austenite grains and effectively refines the grains; preferably 0.015% to 0.025%.
  • Ti 0.01 to 0.03% of Ti is added because the added strong carbide forming element Ti also exhibits a strong effect of suppressing austenite grain growth at a high temperature, and the addition of Ti contributes to effectiveness.
  • the grain is refined; it is preferably controlled at Ti: 0.015% to 0.025%.
  • the invention has a unique refining grain process: a low-temperature coiling process in the bainite phase transformation zone of the hot rolling process, and the process is designed with a Cr-free and Mo-component system.
  • the hot coil tensile strength is less than 1000 MPa, to ensure the cold roll manufacturability; the refinement of Cr and Mo grains and the effect of improving the strength of the material, the rational design of the composition and the process are also achieved in the present invention, and cut costs.
  • the inevitable impurity elements include P, N, and S
  • the lower the impurity content is controlled, the better the implementation effect is, the mass percentage of P is controlled at P ⁇ 0.015%, and the MnS formed by S is severe.
  • the forming property is affected, and thus the mass percentage of S is controlled to be S ⁇ 0.003%, and since N easily causes cracks or bubbles on the surface of the slab, N ⁇ 0.005%.
  • the content of C and Mn alloy elements needs to have upper limit control to ensure excellent welding performance and molding performance, and to avoid the strength exceeding the upper limit (the upper limit of strength of each standard of 1180 MPa high-strength steel for high forming energy is different, and 1350 MPa is a more common standard) 1300 MPa is a stricter standard, and the strength exceeding the upper limit is bound to bring about a decrease in molding performance).
  • the micro-addition of Nb and Ti suppresses the effect of austenite grain growth and effectively refines crystal grains.
  • the C content of the ferrite transformation zone is shifted to the right, and vice versa; the Mn content is increased, and the bainite transformation is shifted upward, and vice versa. Therefore, when the C content is high, the Mn content can be controlled to the lower limit; when the C content is low, the Mn content needs to be controlled to the upper limit.
  • the addition of Si also shifts the ferrite phase transition zone to the right, but the effect is less than the increase in the same C content. Therefore, in the present invention, the C, Mn, and Si contents are also required to conform to the formula: 1.73 ⁇ [C] ⁇ [Mn] + [Si] ⁇ 1.
  • the content of Al and N in the component system of the present invention is integrated, and the addition of Nb and Ti needs to be ensured to a certain amount to be refined.
  • the role of the granules. Therefore, in the present invention, the Nb and Ti contents are also required to conform to the formula: [Nb] + [Ti] ⁇ 3 ⁇ 0.047.
  • the method for manufacturing a low-cost, high-formability 1180 MPa grade cold-rolled annealed duplex steel sheet according to the present invention comprises the steps of:
  • tempering temperature is 200-270 ° C
  • tempering time is 100-400 s
  • hot rolling is first heated to 1100-1230 ° C, and the temperature is maintained for more than 0.6 hours, and then hot rolling is performed at a temperature higher than Ar 3 , and then rolled. Rapid cooling at a rate of 30-100 ° C / s.
  • the special coiling temperature is adopted: the low-temperature coiling (400-500 °C) of the bainite phase transformation zone in the hot rolling process ensures that the bainite content of the hot-rolled structure is 80% or more.
  • the cooling rate is not more than 0.3 ° C / s and is slowly cooled to below 370 ° C, and does not enter the martensite transformation zone, ensuring that the hot rolling strength meets the requirements of cold rolling manufacturability.
  • the annealing soaking temperature is limited to 790-840 ° C, because at this annealing temperature, the tensile strength of 1180 MPa can be ensured, and the obtained grain size can be ensured to be fine, thereby obtaining Good molding performance.
  • Annealing homogenization below 790 °C can not obtain 1180MPa tensile strength; annealing homogenization is higher than 840 °C, the grain size is coarse, and the molding performance is greatly reduced.
  • the soaking temperature of 800-820 ° C, C content of 0.11-0.125% the grain size obtained is finer, and the mechanical properties can be obtained moderately, the performance better.
  • the fine grain size is advantageous for ensuring good elongation and bending properties, so 0.11-0.125% C is a better C content range, and 800-820 ° C is a better continuous annealing soak temperature process window.
  • the present invention obtains a tensile strength of more than 1180 MPa and a fine uniform martensite + ferrite dual phase structure without the addition of Mo and Cr by rational design of the alloy composition and process.
  • Steel plate Under the premise of ensuring high strength, the elongation and cold bending performance are excellent.
  • the yield strength is greater than 850 MPa; the tensile strength is greater than 1180 MPa; the elongation is ⁇ 8%; the 90 degree cold bending performance is characterized by the parameter R/t ⁇ 2.5, where R is the bending radius and t is the thickness of the plate, in mm.
  • the manufacturing method of the present invention obtains a high-formability cold-rolled ultra-high-strength dual-phase steel sheet with tensile strength greater than 1180 MPa, low cost, high elongation, and good cold bending performance through alloy component design and specific process parameter design. .
  • Table 1 The composition of the steel embodiment of the present invention is shown in Table 1, and the balance of the ingredients is Fe.
  • Table 2 lists the process parameters of the steel sheets of the examples.
  • Table 3 lists the relevant performance parameters for the steel sheets of the examples.
  • Hot rolling first heating to 1100-1230 ° C, holding for more than 0.6 hours, then hot rolling with Ar3 or higher, rapid cooling at 30-100 ° C / s after rolling; bainite phase transformation zone in hot rolling process Low temperature coiling (400-500 ° C), to ensure that the final microstructure of hot-rolled bainite content is more than 80%; after cooling, the cooling rate is not more than 0.3 ° C / s and slowly cooled to below 370 ° C, does not enter the martensite transformation zone, Ensure that the hot rolling strength meets the requirements for cold rolling manufacturability;
  • tempering temperature is 200-270 ° C
  • tempering time is 100-400 s
  • Examples 1-18 are low-cost, high-formability 1180 MPa grade cold-rolled annealed duplex steel sheets obtained under the composition and process of the present invention, and obtained mechanical properties: yield strength greater than 850 MPa; tensile strength The strength is greater than 1180 MPa, the elongation is ⁇ 8%, and the 90 degree cold bending performance is characterized by a parameter R/t ⁇ 2.5. Therefore, the low-cost and high-formability 1180 MPa grade cold-rolled annealed duplex steel sheet according to the present invention has a tensile strength of more than 1180 MPa and a good elongation ratio without adding precious alloy elements such as Mo and Cr. Excellent cold bending performance.

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Abstract

一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法，其化学成分质量百分比为：C：0.1％～0.125％，Si：0.4％～0.8％，Mn：2.6％～2.9％，Al：0.01％～0.05％，Nb：0.01～0.03％，Ti：0.01～0.03％，余量为Fe和不可避免杂质。本发明所述双相钢板通过合金元素和制造工艺的合理设计，保证了钢板在低成本前提下达到1180MPa级强度；获得细小均匀的马氏体+铁素体双相组织以保证延伸率和冷弯性能上表现优良，具有较好成型性。所述双相钢板的屈服强度大于850MPa；其抗拉强度大于1180MPa；其延伸率≥8％；90度冷弯性能表征参量R/t≤2.5。

Description

一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法 技术领域
本发明涉及双相钢板及其制造方法，尤其涉及一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法，该钢板具有850MPa以上的屈服强度，1180MPa以上的抗拉强度，优良的成型性能，尤其是优良的冷弯性能。
背景技术
随着全球能源危机和环境问题的加剧，节能和安全成为了汽车制造业的主要发展方向。降低车重是节能和减少排放措施之一。高强度双相钢具有良好的机械性能和使用性能，适于结构件的制造。
随着市场的变化，对于更低成本、更优性能的要求变得日益突出。目前780DP、980DP是应用的主流，1180DP由于成型性较差，主要用于简单成型的结构件。随着减重节能的趋势不断发展，未来双相钢的中心强度级别预计会上移，1180DP的需求将增加，但前提是要求1180DP具有和980DP相当的延伸率和弯曲性能。
现有冷轧超高强度双相钢如中国专利申请公开号CN102227511B公开了一种具有1180MPa以上的TS、且扩孔性和弯曲性等成形性优良的高强度冷轧钢板、高强度热镀锌钢板及它们的制造方法；其C：0.05-0.3％，Si：0.5-2.5％，Mn：1.5-3.5％，P：0.001-0.05％，S：0.0001-0.01％，Al：0.001-0.1％，N：0.0005-0.01％，Cr1.5％以下且包括0％。要求含有铁素体相和马氏体相，所述马氏体相在组织整体中所占的面积率为30％以上，(马氏体相所占的面积)/(铁素体相所占的面积)大于0.45且小于1.5。其独特性在于依靠高比例的铁素体含量获得较好成型性。
中国专利申请公开号CN105829564A提供一种适用于汽车部件用坯材的、屈服强度(YS)为780MPa以上、拉伸强度(TS)为1180MPa以上且点焊性、延展性和弯曲加工性优异的高强度钢板及其制造方法；其C： 0.05-0.15％，Si：0.01-1.00％，Mn：1.5-4.0％，P：0.1％以下，S：0.02％以下，Al：0.01-0.5％，Nb：0.005-0.1％，Ti：0.005-0.1％，Cr：0.01-2.0％，B:0.0005-0.005％。要求将铁素体和马氏体的平均结晶粒径设为10μm以下。其独特性在于依靠B元素的添加及相应工艺细化晶粒和提高淬透性，且添加Cr提高淬透性。
中国专利申请公开号CN102828106A公开的钢板，抗拉强度在1180MPa以上，加工性和低温脆性优异。本发明的高强度钢板，含有C：0.10～0.30％、Si：1.40～3.0％、Mn：0.5～3.0％、P：0.1％以下、S：0.05％以下、Al：0.005～0.20％、N：0.01％以下、O：0.01％以下，余量由Fe和不可避免的杂质构成，并且，相对于全部组织，初生马氏体和残留奥氏体的混合组织(MA组织)的体积率为6％以下(不含0％)，并且以X射线衍射法测量残留奥氏体时，相对于全部组织的残留奥氏体的体积率为5％以上。其独特性在于利用Si：1.40～3.0％的添加，在最终组织获得残余奥氏体，提高成型性能。
中国专利申请公开号CN106661701A种用于通过对钢板进行热处理来生产高强度钢板的方法，所述高强度钢板的屈服强度YS>850MPa，拉伸强度TS>1180MPa，总延伸率>13％，且扩孔率HER>30％，其中所述钢的化学组成包含：0.13％≤C≤0.22％，1.2％≤Si≤1.8％，1.8％≤Mn≤2.2％，0.10％≤Mo≤0.20％，Nb≤0.05％，Ti<0.05％，Al≤0.5％，剩余部分为Fe和不可避免的杂质。最终组织能够包含3％至15％的残余奥氏体。其独特性在于利用1.2％≤Si≤1.8％的添加，在最终组织获得残余奥氏体，提高成型性能。
中国专利申请公开号CN106661653A用于制造屈服强度YS大于1000MPa，拉伸强度TS大于1150MPa且总延伸率E大于8％的钢板的方法，所述方法包括以下步骤：通过轧制钢来制备钢板，以重量百分比计，所述钢包含0.19％至0.22％的C，2％至2.6％的Mn，1.45％至1.55％的Si，0.15％至0.4％的Cr，小于0.020％的P，小于0.05％的S，小于0.08％的N，0.015％至0.070％的Al，在115秒至240秒的时间期间内将钢板加热至高于380℃的第一过时效温度TOA1，然后在300秒至610秒的时间期间内将板加热到在420℃与450℃之间的第二过时效温度TOA2，使钢板以小 于5℃/秒的冷却速度冷却至低于100℃的温度。钢的组织包含大于80％的回火马氏体，大于5％的残余奥氏体。其独特性在于利用0.19％至0.22％的高C含量、1.45％至1.55％的Si含量，配合独特2次时效的工艺获得残余奥氏体加马氏体组织，保证高强度和高成型性。
发明内容
本发明的目的在于提供一种低成本高成型性1180MPa级冷轧退火双相钢板及其制造方法，通过合金元素和制造工艺的合理设计，保证钢板在低成本前提下达到1180MPa级强度；获得细小均匀的马氏体+铁素体双相组织以保证延伸率和冷弯性能上表现优良，具有较好成型性；所述双相钢板的屈服强度大于850MPa，其抗拉强度大于1180MPa，其延伸率≥8％，90度冷弯性能表征参量R/t≤2.5，其中R表示弯曲半径，t表示板厚，单位mm。
为达到所示目的，本发明的技术方案是：
本发明钢成分设计以C+Mn为主的成分体系，保证达到1180MPa级强度，该钢不添加Mo、Cr等贵重合金元素保证低成本。Nb、Ti的微量添加，达到抑制奥氏体晶粒长大的效果，有效细化晶粒。不添加Mo、Cr的特殊成分设计和特殊的热轧低温卷取工艺配合，保证热卷抗拉强度小于1000MPa，保证冷轧的可制造性。
本发明制造工艺方面，由于不添加微合金元素Cr、Mo不利细化晶粒，所以采用特殊的热轧卷取温度：热轧工序贝氏体相变区低温卷取(400-500℃)，保证热轧组织贝氏体含量80％以上；卷取后冷速不大于0.3℃/s至370℃以下，不进入马氏体相变区域，保证热轧强度满足冷轧可制造性要求。通过以上工艺，冷轧态组织为均匀的具有大量位错等缺陷的变形贝氏体组织，提供了大量的奥氏体形核点，使得奥氏体呈现爆发式的形核，因此奥氏体晶粒明显细化。且这些缺陷也成为了碳原子高速扩散的通道，又因为整个基体组织统一，碳浓度差小，使得奥氏体快速生成，体积分数增大得到细小均匀最终组织。
具体的，本发明的一种低成本高成型性1180MPa级冷轧退火双相钢板，其化学成分质量百分比为：C：0.1％～0.125％，Si：0.4％～0.8％， Mn：2.6％～2.9％，Al：0.01％～0.05％，Nb：0.01～0.03％，Ti：0.01～0.03％，余量为Fe和不可避免杂质。
本发明所述双相钢板的组织形态为细小均匀的马氏体+铁素体双相组织，其中，含有至少75％以上的马氏体，其余为铁素体；马氏体的晶粒直径不大于5微米，铁素体的晶粒直径不大于5微米。
本发明所述双相钢板的屈服强度大于850MPa；其抗拉强度大于1180MPa；其延伸率≥8％；90度冷弯性能表征参量R/t≤2.5，其中R表示弯曲半径，t表示板厚，单位mm。
在本发明所述双相钢板的设计中：
C：在本发明所述的双相钢板中，C元素的添加起到提高钢的强度，提高马氏体的硬度。因此，选择C的质量百分比在0.1％～0.125％之间，这是因为：当C的质量百分比低于0.1％，则钢板的强度受到影响，并且不利于奥氏体的形成量和稳定性；当C的质量百分比高于0.125％，则造成马氏体硬度过高，晶粒尺寸粗大，不利于钢板的成型性能。优选为0.11％～0.125％。
Si：添加Si可以提高淬透性。并且钢中固溶的Si可以影响位错的交互作用，增加加工硬化率，在双相钢中可以适当提高延伸率，有益于获得较好的成型性。Si含量控制在Si：0.4％～0.8％，优选为0.5％～0.7％。
Mn：添加Mn元素有利于钢的淬透性提高，有效提高钢板的强度。而选取Mn的质量百分比在2.6％～2.9％是因为：当Mn的质量百分比低于2.6％时，则钢板的强度不足；当Mn的质量百分比高于2.9％时，则钢板的强度过高，使得其成型性能下降，且容易发生偏析。因此，本发明所述的低成本高成型性1180MPa级冷轧退火双相钢板中控制Mn的质量百分比在Mn：2.6-2.9％，优选为2.7％～2.85％。
Al：添加Al起到了脱氧作用和细化晶粒的作用，因此，Al的质量百分比控制在Al：0.01％～0.05％，优选为0.015～0.045％。
Nb：添加0.01～0.03％的Nb，是因为：在微合金钢中加入少量的强碳化物形成元素Nb后，在控制轧制过程中，应变诱导析出相通过质点钉扎和亚晶界的作用而相当显著的降低变形奥氏体的再结晶温度，提供形核质点，对细化晶粒作用明显；在连退奥氏体化过程中，均热未溶的碳、氮 化物质点将通过质点钉扎晶界机制而阻止均热奥氏体晶粒的粗化，有效细化晶粒；优选为0.015％～0.025％。
Ti：添加0.01～0.03％的Ti，是因为：添加的强碳化物形成元素Ti在高温下也显示出一种强烈的抑制奥氏体晶粒长大的效果，同时Ti的添加有助于有效细化晶粒；优选控制在Ti：0.015％～0.025％。
Cr、Mo：不添加Cr、Mo。首先，此设计可降低钢的成本；第二，本发明具有独特的细化晶粒工艺：热轧工序贝氏体相变区间低温卷取的工艺，此工艺是配合无Cr、Mo成分体系设计的，这样才能保证热卷抗拉强度小于1000MPa，保证冷轧可制造性；Cr、Mo的细化晶粒和提高材料强度的作用，本发明中通过成分和工艺的合理设计同样做到，且降低成本。
在本发明所述的技术方案中，不可避免的杂质元素包括P、N、S，杂质含量控制得越低，实施效果越好，P的质量百分比控制在P≤0.015％，S形成的MnS严重影响成形性能，因而S的质量百分比控制在S≤0.003％，由于N容易导致板坯表面产生裂纹或气泡，因而，N≤0.005％。
在上述成分设计中，不添加Mo、Cr等贵重合金元素保证低成本，同时为了保证在30-100℃/s正常连续退火气体冷却冷却速度下获得1180MPa级抗拉强度，成分中需要保证C、Mn的合金添加含量以提供足够的淬透性。但C、Mn合金元素的含量需要有上限控制，以保证优良的焊接性能及成型性能，避免强度超出上限(对高成型能1180MPa级高强钢各标准中强度上限不一，1350MPa为较普遍的标准，1300MPa为较严格的标准，强度超过上限势必带来成型性能的下降)。Nb、Ti的微量添加，抑制奥氏体晶粒长大的效果，有效细化晶粒。
另外，为配合热轧工序冷却进入贝氏体相变区过程中不触碰铁素体相变区，保证热轧组织贝氏体含量80％以上，需要控制C含量和Mn含量的配比。本发明所述成分体系相图中：C含量升高铁素体相变区右移，反之左移；Mn含量升高贝氏体相变上移，反之下移。所以C含量较高时，Mn含量可向下限控制；C含量较低时，Mn含量需向上限控制。添加Si也可使铁素体相变区右移，但效果小于相同C含量增加。所以本发明中C、Mn、Si含量还需符合公式：1.73×[C]×[Mn]+[Si]≥1。
由于钢生产过程中Al的氮化物与Nb、Ti的碳氮化物的竞争析出关 系，综合本发明成分体系中Al、N的含量，Nb、Ti的添加需要保证一定的量才能起到细化晶粒的作用。所以本发明中Nb、Ti含量还需符合公式：[Nb]+[Ti]×3≥0.047。
本发明所述的低成本高成型性1180MPa级冷轧退火双相钢板的制造方法，其步骤包括：
1)冶炼、铸造，按上述成分冶炼、铸造成坯；
2)热轧，先加热至1100-1230℃，保温0.6小时以上，然后采用Ar3以上温度热轧，轧后以30-100℃/s的速度快速冷却；卷取温度：400-500℃，卷取后冷速不大于0.3℃/s缓慢冷却至370℃以下；
3)冷轧，控制冷轧压下率为50-70％；
4)退火，退火均热温度为790-840℃，优选800-820℃；然后以v＝3-20℃/s的速度冷却到快冷开始温度T，其中快冷开始温度T≥800-30×v，然后再以30-100℃/s的速度冷却到200-270℃；
5)回火，回火温度为200-270℃，回火时间为100-400s；
6)平整，采用0-0.3％的平整。
在本发明制造方法中，
在本发明所述的低成本高成型性1180MPa级冷轧退火双相钢板的制造方法中，热轧先加热至1100-1230℃，保温0.6小时以上，然后采用Ar3以上温度热轧，轧后以30-100℃/s的速度快速冷却。采用特殊的卷取温度：热轧工序贝氏体相变区低温卷取(400-500℃)，保证热轧组织贝氏体含量80％以上。卷取后冷速不大于0.3℃/s缓慢冷却至370℃以下，不进入马氏体相变区域，保证热轧强度满足冷轧可制造性要求。
在所述退火步骤中，退火均热温度限定为790-840℃，是因为，在该退火温度下，即能保证获得1180MPa的抗拉强度，又能保证获得的晶粒尺寸细小，从而得到较好的成型性能。退火均热低于790℃无法获得1180MPa抗拉强度；退火均热高于840℃，晶粒尺寸粗大，成型性能大幅下降。
综合对比各成分在相同均热温度下的组织，均热温度800-820℃时，C含量在0.11-0.125％时，获得的晶粒尺寸更加细小，且能获得的机械性能较适中，成性能更优。细小的晶粒尺寸有利于保证良好的延伸率和弯曲 性能，所以0.11-0.125％C是更优的C含量范围，800-820℃是更优的连续退火均热温度工艺窗口。
相较于现有技术，本发明通过合金成分和工艺的合理设计，在不添加Mo、Cr的前提下，获得抗拉强度大于1180MPa且兼具细小均匀马氏体+铁素体双相组织的钢板。在保证高强度的前提下，延伸率和冷弯性能上表现优良。其屈服强度大于850MPa；抗拉强度大于1180MPa；其延伸率≥8％；90度冷弯性能表征参量R/t≤2.5，其中R表示弯曲半径，t表示板厚，单位mm。
相应地，本发明所述的制造方法通过合金成分设计及具体工艺参数设计，获得抗拉强度大于1180MPa，低成本，延伸率高，冷弯性能佳的高成型性冷轧超高强度双相钢板。
具体实施方式
下面将结合具体的实施例对本发明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。
本发明钢实施例的成分参见表1，其成分余量为Fe。表2列出了实施例钢板的工艺参数。表3列出了实施例钢板的相关性能参数。
本发明钢实施例的制造方法如下：
(1)冶炼和铸造：获得要求的合金成分，尽量降低S、P的含量。
(2)热轧，先加热至1100-1230℃，保温0.6小时以上，然后采用Ar3以上温度热轧，轧后以30-100℃/s的速度快速冷却；热轧工序贝氏体相变区低温卷取(400-500℃)，保证热轧最终组织贝氏体含量80％以上；卷取后冷速不大于0.3℃/s缓慢冷却至370℃以下，不进入马氏体相变区域，保证热轧强度满足冷轧可制造性要求；
(3)冷轧，控制冷轧压下率为50-70％；
(4)退火，退火均热温度为790-840℃，优选800-820℃，然后以v＝3-20℃/s的速度冷却到快冷开始温度T，其中快冷开始温度T≥800-30×v，然后再以30-100℃/s的速度冷却到200-270℃；
(5)回火，回火温度为200-270℃，回火时间为100-400s；
(6)平整，采用0-0.3％的平整；
从表3可以看出，实施例1-18为本发明所述成分和工艺下获得的低成本高成型性1180MPa级冷轧退火双相钢板，获得的机械性能：其屈服强度大于850MPa；抗拉强度大于1180MPa，延伸率≥8％，90度冷弯性能表征参量R/t≤2.5。由此说明本发明所述的低成本高成型性1180MPa级冷轧退火双相钢板在未添加Mo、Cr等贵重合金元素的前提下，获得了大于1180MPa的抗拉强度，且延伸率较好，冷弯性能优良。
Figure PCTCN2018092876-appb-000001
Figure PCTCN2018092876-appb-000002
Figure PCTCN2018092876-appb-000003
Figure PCTCN2018092876-appb-000004

Claims (10)

  1. 一种低成本高成型性1180MPa级冷轧退火双相钢板，其化学成分质量百分比为：C：0.1％～0.125％，Si：0.4％～0.8％，Mn：2.6％～2.9％，Al：0.01％～0.05％，Nb：0.01～0.03％，Ti：0.01～0.03％，余量为Fe和其他不可避免杂质；且，满足1.73×[C]×[Mn]+[Si]≥1，[Nb]+[Ti]×3≥0.047。
  2. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述C含量优选为0.11％～0.125％。
  3. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述Si含量优选为0.5％～0.7％。
  4. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述Mn含量优选为2.7％～2.85％。
  5. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述Al含量优选为0.015～0.045％。
  6. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述Nb含量优选为0.015％～0.025％。
  7. 如权利要求1所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述Ti含量优选为0.015％～0.025％。
  8. 如权利要求1～7任何一项所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述双相钢板的组织形态为细小均匀的马氏体+铁素体双相组织，其中，含有至少75％以上的马氏体，其余为铁素体；马氏体的晶粒直径不大于5微米，铁素体的晶粒直径不大于5微米。
  9. 如权利要求1～8任何一项所述的低成本高成型性1180MPa级冷轧退火双相钢板，其特征在于：所述双相钢板的屈服强度大于850MPa；其抗拉强度大于1180MPa；其延伸率≥8％；90度冷弯性能表征参量R/t≤2.5，其中R表示弯曲半径，t表示板厚，单位mm。
  10. 如权利要求1～9任何一项所述的低成本高成型性1180MPa级冷轧退火双相钢板的制造方法，其特征在于，包括步骤：
    1)冶炼、铸造，按权利要求1～9任何一项所述的成分冶炼、铸造成 坯；
    2)热轧，先加热至1100-1230℃，保温0.6小时以上，然后采用Ar3以上温度热轧，轧后以30-100℃/s的速度快速冷却；卷取温度：400-500℃，卷取后冷速不大于0.3℃/s缓慢冷却至370℃以下；
    3)冷轧，控制冷轧压下率为50-70％；
    4)退火，退火均热温度为790-840℃，优选800-820℃；然后以v＝3-20℃/s的速度冷却到快冷开始温度T，其中快冷开始温度T≥800-30×v，然后再以30-100℃/s的速度冷却到200-270℃；
    5)回火，回火温度为200-270℃，回火时间为100-400s；
    6)平整，采用0-0.3％的平整。
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