WO2025002398A1 - 一种高r值的超高强钢板及其制造方法 - Google Patents

一种高r值的超高强钢板及其制造方法 Download PDF

Info

Publication number
WO2025002398A1
WO2025002398A1 PCT/CN2024/102541 CN2024102541W WO2025002398A1 WO 2025002398 A1 WO2025002398 A1 WO 2025002398A1 CN 2024102541 W CN2024102541 W CN 2024102541W WO 2025002398 A1 WO2025002398 A1 WO 2025002398A1
Authority
WO
WIPO (PCT)
Prior art keywords
steel plate
surface layer
value
strength steel
high strength
Prior art date
Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
Ceased
Application number
PCT/CN2024/102541
Other languages
English (en)
French (fr)
Inventor
朱晓东
薛鹏
李伟
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
Baoshan Iron and Steel Co Ltd
Original Assignee
Baoshan Iron and Steel Co Ltd
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Baoshan Iron and Steel Co Ltd filed Critical Baoshan Iron and Steel Co Ltd
Priority to KR1020267002712A priority Critical patent/KR20260032574A/ko
Priority to EP24831025.2A priority patent/EP4737603A1/en
Publication of WO2025002398A1 publication Critical patent/WO2025002398A1/zh
Anticipated expiration legal-status Critical
Ceased legal-status Critical Current

Links

Classifications

    • 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
    • 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
    • C21D3/00—Diffusion processes for extraction of non-metals; Furnaces therefor
    • C21D3/02—Extraction of non-metals
    • C21D3/04—Decarburising
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties of ferrous metals or ferrous alloys 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
    • 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 of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties of ferrous metals or ferrous alloys by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C33/00—Making ferrous alloys
    • C22C33/04—Making ferrous alloys by melting
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • 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
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • 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 present invention relates to a steel plate and a manufacturing method thereof, and in particular to a high-strength steel plate and a manufacturing method thereof.
  • the r value represents the ability of the steel plate to resist thinning during the deformation process of drawing.
  • the r value is one of the important technical indicators.
  • a high r value means that the steel plate has better drawing performance.
  • the Chinese patent document with the publication number CN101768695A and the publication date of July 7, 2010, entitled "1000MPa Ti microalloyed ultrafine grain cold-rolled dual-phase steel and its preparation process” discloses that its chemical composition is: 0.03-0.2% C, 0.2-0.8% Si, 1.2-2.0% Mn, Ti: 0.03-0.15%, ⁇ 0.02% P, S ⁇ 0.015, 0.02-0.15% Al, and the rest is Fe and unavoidable impurities.
  • annealing in the critical zone with a cooling rate of less than 50°C/s is performed to obtain a cold-rolled dual-phase steel with a strength of more than 980MPa. It can be seen that the high-strength steel does not pay attention to the r value of the steel plate.
  • the Chinese patent document with publication number CN101363099A, publication date February 11, 2009, and titled "A cold-rolled dual-phase steel plate with a tensile strength of 1000 MPa and a preparation method thereof" discloses that C: 0.14-0.21%, Si: 0.4-0.9%, Mn: 1.5-2.1%, P: ⁇ 0.02%, S ⁇ 0.01%, Nb: 0.001-0.05%, V: 0.001-0.02%, after hot rolling and cold rolling, it is kept at 760-820°C, the cooling rate is 40-50°C/s, and it is over-aged at 240-320°C for 180-300s.
  • This high-strength steel does not pay attention to the r value of the steel plate.
  • Xiong Ziliu et al. introduced a dual-phase steel with a tensile strength of 980 MPa in "Forming Properties of High-Strength Dual-Phase Steel” (Metal Heat Treatment, Vol. 46, No. 5, 2021), but its r value is lower than 0.9.
  • Ge Delong et al. introduced 1000-1400MPa martensitic steels in "Maximum Bending Angle of 1000MPa Ultra-High Strength Martensitic Steel Based on Three-Point Bending" (Journal of Plasticity Engineering, Vol. 21, No. 4, 2014), whose r value is less than 0.7.
  • One of the purposes of the present invention is to provide a high r-value ultra-high strength steel plate, which has ultra-high strength and high r-value at the same time.
  • the high r-value ultra-high strength steel plate has a lighter thickness reduction during deformation, thereby showing better formability regardless of global tensile deformation or local deformation.
  • the r-value represents the ability of the steel plate to resist thinning during the deformation process of drawing.
  • the thickness of the steel plate is not easy to be thinned during the tensile deformation, so it is not easy to break.
  • the r-value when it undergoes severe deformation locally, if the r-value is high, the thickness is not easy to be thinned locally, which can slow down the necking process, thus being beneficial to the formability of the steel plate.
  • the present invention proposes a high r-value ultra-high strength steel plate, which includes an upper surface layer, an intermediate layer and a lower surface layer in the thickness direction; wherein the main body of the microstructure of the upper surface layer and the lower surface layer is ferrite, the carbon content of the upper surface layer and the lower surface layer is ⁇ 0.025%, and the mass percentage of the chemical elements in the upper surface layer and the lower surface layer respectively satisfies: Ti-3.42N-3.98C ⁇ 0; the microstructure of the intermediate layer includes ferrite and tempered martensite.
  • One of the cores of the present invention is that a high r value is achieved by making the mass percentage of chemical elements in the upper and lower surface layers respectively satisfy Ti-3.42N-3.98C ⁇ 0, and the carbon content of the upper and lower surface layers is ⁇ 0.025%. At the same time, the high strength of the steel plate as a whole is mainly guaranteed by the organization of the middle layer.
  • the carbon content of the upper and lower surface layers is 0.02-0.025%.
  • the mass percentage of chemical elements in the upper and lower surface layers respectively satisfy: Ti-3.42N-3.98C ⁇ 0.060, preferably ⁇ 0.068. In some embodiments, the mass percentage of chemical elements in the upper and lower surface layers respectively satisfy: 0.16 ⁇ Ti-3.42N-3.98C ⁇ 0.06.
  • the average grain diameter of the ferrite in the upper surface layer and the lower surface layer is ⁇ 15 microns. In some embodiments, in the high r value ultra-high strength steel plate of the present invention, the average grain diameter of the ferrite in the upper surface layer and the lower surface layer may be ⁇ 10 microns. In some embodiments, in the high r value ultra-high strength steel plate of the present invention, the upper surface layer and the lower surface layer may be ⁇ 10 microns. The average grain diameter of ferrite in the surface layer and the lower surface layer is 6-10 microns.
  • the volume phase ratio of ferrite in the upper surface layer and the lower surface layer is ⁇ 97% respectively. In some embodiments, the volume phase ratio of ferrite in the upper surface layer and the lower surface layer is 97-99%.
  • the microstructure of the intermediate layer may also include bainite.
  • the volume phase ratio of tempered martensite + bainite is ⁇ 40%.
  • the intermediate layer contains ferrite, and at least one of tempered martensite and bainite.
  • the volume phase ratio of ferrite in the intermediate layer is 55-60%, preferably 57-60%.
  • the volume phase ratio of tempered martensite in the intermediate layer is 40-45%, preferably 40-43%.
  • the volume phase ratio of bainite in the intermediate layer is 0-5%, preferably 0-3%.
  • the ferrite of the upper surface layer and the lower surface layer contains carbides, and the types of the carbides are Ti(C, N) and Nb(C, N).
  • the thickness of the upper surface layer and the lower surface layer are respectively 100-200 microns.
  • the sum of the thickness of the upper surface layer and the lower surface layer does not exceed 40% of the total thickness of the steel plate.
  • the thickness of the high r-value ultra-high strength steel plate of the present invention is 0.8-2.5 mm, preferably 1.2-2.0 mm.
  • the high r-value ultra-high strength steel plate of the present invention contains Fe and inevitable impurity elements, and further contains the following chemical elements in the following mass percentages:
  • the mass percentage of each chemical element is:
  • the design principles of the chemical elements of the high r-value ultra-high strength steel plate of the present invention are specifically as follows:
  • C In the high r-value high-strength steel plate described in the present invention, C can improve the strength of the steel by affecting the hardness of martensite. If the carbon content in the steel is too low, the strength of the steel after quenching is low; and the higher the carbon content in the steel, the higher the strength of the steel after quenching. However, too high carbon is not conducive to obtaining a high r-value. Therefore, in the high r-value ultra-high strength steel plate described in the present invention, the C element content is controlled between 0.08-0.20%.
  • Si In the high-r value high-strength steel plate of the present invention, Si can play the role of solid solution strengthening, inhibiting the precipitation of Fe3C, and promoting the formation of retained austenite. Based on this, the present invention controls the Si content to be 0.01-1.7%. In some embodiments, the Si content is 0.4-1.7%.
  • Mn is the main element for improving the hardenability of steel.
  • the Mn content matches the cooling capacity of the cooling method used. When the Mn content is too low, the steel plate cannot be hardened, and thus cannot obtain a high strength. If the Mn content is too high, it is not good for the carbon equivalent. Therefore, in the high r value ultra-high strength steel plate described in the present invention, the Mn element content is controlled between 0.7-2.7%. In some embodiments, the Mn content is 1.0-2.7%.
  • Ti In the high r value high strength steel plate of the present invention, the content of Ti is closely related to the r value of the present invention. If the Ti content is too low, it is insufficient to fix all the C and N in the upper and lower surface layers, resulting in a low r value; if the Ti content is too high, Ti is excessive. Therefore, in the high r value ultra-high strength steel plate of the present invention, the Ti element is controlled between 0.16-0.26%.
  • Nb plays a role in assisting the fixation of C and N, forming Nb (C, N), and ensuring that C and N in the upper and lower surface regions are not in a solid solution state. Therefore, in the high r value ultra-high strength steel plate of the present invention, the Nb element is controlled between 0.015-0.05%.
  • Al Its main function in the present invention is deoxidation during the smelting process, and its content range can be: 0.02-0.06%.
  • N is not conducive to forming a composition ratio of Ti-3.42N-3.98C ⁇ 0 in the upper and lower surface layers, and is not conducive to forming a state of no interstitial atoms locally, so it is necessary to control N ⁇ 0.005%.
  • the inevitable impurities are mainly S and P.
  • the content of P element is ⁇ 0.015%.
  • the content of S element is ⁇ 0.003%.
  • the basic components are: C: 0.08-0.20%; Si: 0.01-1.7%; Mn: 0.7-2.7%.
  • This basic component constitutes the basis of the strength of the steel plate, but a high r value cannot be obtained based on these basic components.
  • the steel plate described in the present invention also contains: Ti: 0.16-0.26%, Nb: 0.015-0.05%, N ⁇ 0.005%.
  • the chemical elements also contain at least one of Cr, Mo, and B; wherein B ⁇ 0.005%, Cr+Mo ⁇ 0.5%.
  • the content of B does not exceed 0.005%
  • the content of Cr does not exceed 0.3%
  • the content of Mo does not exceed 0.2%.
  • the alloy elements Cr, Mo, and B can effectively cooperate with manganese to improve the hardenability, so as to match the hardenability and rapid cooling capacity of the steel, so as to further avoid failure to harden or excessively high carbon equivalent.
  • the tensile strength of the high r value ultra-high strength steel plate of the present invention is ⁇ 980MPa, its r value r 90 ⁇ 1.2, the microhardness of the upper surface layer and the lower surface layer HV ⁇ 130, and the microhardness of the middle layer HV ⁇ 330.
  • the tensile strength of the high r value ultra-high strength steel plate of the present invention is ⁇ 1020MPa.
  • the tensile strength of the high r value ultra-high strength steel plate of the present invention may be 980-1100MPa, preferably 980-1060MPa.
  • the high r-value ultra-high strength steel plate of the present invention has an r-value R90 of 1.2-1.4.
  • the microhardness HV of the upper surface layer and the lower surface layer of the high r-value ultra-high strength steel plate of the present invention is independently 100-130.
  • the microhardness HV of the middle layer of the high r-value ultra-high strength steel plate of the present invention is 330-380.
  • the elongation of the high r-value ultra-high strength steel plate of the present invention is ⁇ 12%, such as 12-16% or 13-15%.
  • the yield strength of the high r-value ultra-high strength steel plate of the present invention is ⁇ 570 MPa, such as 570-660 MPa.
  • another object of the present invention is to provide a method for manufacturing a high r-value ultra-high strength steel plate, which can produce a high r-value ultra-high strength steel plate by optimizing the process.
  • the present invention also proposes a method for manufacturing a high r-value ultra-high strength steel plate as described above, comprising the steps of: smelting and casting; hot rolling; cold rolling after pickling; annealing; wherein the annealing step comprises:
  • High temperature humidification decarburization process spray water vapor to humidify the annealing furnace, and at the same time control the steel plate heating temperature to 880-950°C for 100-300s, and control the dew point in the annealing furnace to be above 0°C;
  • the steel plate is cooled from the high temperature and humidification temperature to between 730-770°C at a cooling rate of 3-10°C/s;
  • Quasi-static slow cooling process the steel plate is slowly cooled from 730-770°C to 680-720°C, the cooling temperature range is controlled at 40-60°C, the cooling rate is 0.03-0.1°C/s, and the slow cooling process is ensured to last for more than 500s;
  • Rapid cooling process The steel plate is rapidly cooled to below 300°C at a cooling rate of ⁇ 50°C/s;
  • the steel plate is tempered at 240-320°C for 150-500s.
  • the process used in the manufacturing method of the present invention is crucial for the ultra-high strength steel plate of the present invention to obtain a high r-value.
  • the heating temperature of the steel plate is between 880 and 950°C and maintained for 100-300s.
  • the amount of humidified water is controlled according to the weight of the steel strip passing through the unit time and the thickness of the required decarburization layer to ensure that the dew point in the annealing furnace is above 0°C. In some embodiments, the dew point in the annealing furnace is 0-20°C.
  • the high-temperature humidification decarburization process makes the upper and lower surfaces of the steel plate respectively become preliminary decarburization layers with a thickness of 100-200 microns (that is, both the upper and lower surfaces are decarburization layers).
  • the heating temperature of the steel plate is controlled to be 880-950°C, which is conducive to the decarburization reaction of the steel plate.
  • the reason for selecting the dew point in the furnace to be above 0°C is that if the dew point in the furnace is low, the decarburization reaction of the steel plate will be slow, and it is not enough to obtain a decarburization layer of sufficient thickness.
  • the carbon content of the upper and lower surfaces of the steel plate is significantly reduced, which can be reduced to a level of ⁇ 0.05%.
  • the carbon content of the upper and lower surfaces of the steel plate is still too high, and the austenite is directly cooled and transformed into ferrite, and Fe3C will still precipitate.
  • a quasi-static slow cooling process is further carried out below the Ac1 temperature of the iron-carbon phase diagram. In this quasi-static slow cooling process, the steel plate is slowly cooled from 730-770°C to 680-720°C at a rate of 0.03-0.1°C/s.
  • the temperature range of the entire slow cooling process is controlled at 40-60°C, and the duration is ⁇ 500s. In some embodiments, the slow cooling duration is 500-1500s.
  • the upper and lower surfaces of the steel plate that has undergone preliminary decarburization are first subjected to austenite decomposition to form ferrite.
  • the carbon in the ferrite diffuses into the austenite in the high-carbon zone (middle layer) that has not undergone transformation, thereby further causing the carbon content of the upper and lower surfaces to decrease. If the slow cooling time is sufficient, the carbon content of the upper and lower surfaces can be reduced to a level close to that of industrial pure iron, so that the upper and lower surface areas of the invention steel meet Ti-3.42N-3.98C ⁇ 0.
  • slow cooling from 730-770°C to 680-720°C is crucial.
  • the austenite decarburization content of the surface layer decreases, and Ac3 increases, so that the austenite ⁇ ferrite transformation occurs first in the process of slow cooling from 730-770°C to 680-720°C.
  • the local carbon content of the layer is reduced to ⁇ 0.025%, the Ti, C, and N contents in the steel plate can meet the level of Ti-3.42N-3.98C ⁇ 0, thereby essentially forming a structure without gap atoms in the upper and lower surface layers, greatly improving the r value of the upper and lower surface layers, and thereby improving the r value of the steel plate as a whole.
  • the steel plate is rapidly cooled to below 300°C at a cooling rate of ⁇ 50°C/s, and tempered at 240-320°C for 150-500s.
  • the cooling rate of the rapid cooling is 50-150°C/s.
  • the termination temperature of the rapid cooling is 240-300°C.
  • the austenite in the middle layer is transformed into a martensite structure, and in some embodiments, a small amount of bainite and residual austenite structure may also be contained, thereby achieving ultra-high strength, such as a tensile strength of ⁇ 980MPa.
  • the overaging temperature is 240-300°C.
  • the smelting, casting, hot rolling and cold rolling processes in the manufacturing method of the present invention can all adopt conventional processes, and the present invention does not make any special improvements to these processes.
  • the steel plate is heated out of the furnace at a temperature of 1220-1280°C, the final rolling temperature is 870-930°C, and the coiling temperature is 570-630°C.
  • the cold rolling reduction rate is controlled to be 40-65%.
  • the present invention ensures ultra-high strength by controlling the organizational morphology of the middle layer, and at the same time achieves a high r value by controlling the mass percentage of chemical elements in the upper and lower surface layers to respectively satisfy Ti-3.42N-3.98C ⁇ 0, thereby achieving an ultra-high strength steel plate with a high r value.
  • the present invention can substantially achieve a solid solution state of interstitial atoms in the upper and lower surface layers by adding an appropriate amount of Ti and assisting with Nb to fix N and C, thereby greatly improving the r value of the surface layer, which in turn affects the r value of the entire steel plate, so that the r value of the 980MPa cold-rolled high-strength steel plate reaches above 1.2, thereby greatly reducing the thinning rate during the deformation process, thereby obtaining better formability.
  • the high-r-value high-strength steel plate described in the present invention adopts a unique annealing process, undergoes a preliminary high-temperature humidification decarburization process and a long-term slow cooling austenite ⁇ ferrite transformation, so that the carbon content is redistributed, thereby reducing the carbon content of the upper and lower surface layers of the ultra-high-strength steel plate to below 0.025%.
  • FIG. 1 schematically shows a method for manufacturing a high-r value high-strength steel plate according to the present invention in a Annealing process curve under implementation mode.
  • the high-r-value high-strength steel plates of Examples 1-6 and the comparative steel plates of Comparative Examples 1-2 were prepared by the following steps:
  • Hot rolling conventional hot rolling process is adopted. Specifically, the steel plate heating furnace temperature is controlled to be 1220-1280°C, the final rolling temperature is 870-930°C, and the coiling temperature is 570-630°C;
  • High temperature humidification decarburization process spray water vapor to humidify the annealing furnace, and at the same time control the steel plate heating temperature to 880-950°C for 100-300s, and control the dew point in the annealing furnace to be above 0°C;
  • the steel plate is cooled from the high temperature and humidification temperature to between 730-770°C at a cooling rate of 3-10°C/s;
  • Quasi-static slow cooling process the steel plate is slowly cooled from 730-770°C to 680-720°C, the cooling temperature range is controlled at 40-60°C, the cooling rate is 0.03-0.1°C/s, and the slow cooling process is ensured to last for more than 500s;
  • Rapid cooling process The steel plate is rapidly cooled to below 300°C at a cooling rate of ⁇ 50°C/s;
  • the steel plate is tempered at 240-320°C for 150-500s.
  • Table 1 lists the mass percentages of the chemical elements of Examples 1-6 of the present invention and Comparative Examples 1-2.
  • Table 2-1 lists the specific process parameters of the annealing step of the high r-value high strength steel plates of Examples 1-6 and the comparative steel plates of Comparative Examples 1-2.
  • Ti-3.42N-3.98C value in the upper and lower surface layers is greater than 0, it indicates that there are no interstitial atoms in the upper and lower surface layers, and when the Ti-3.42N-3.98C value is less than 0, it indicates that there are interstitial atoms in the upper and lower surface layers.
  • the two columns of data in the table represent the values of the upper and lower surface layers respectively.
  • tensile test is carried out along the direction perpendicular to the rolling direction of the steel plate.
  • JIS5# standard plate tensile test specimen with a gauge length of 50mm and a gauge area width of 25mm, the test is carried out on a tensile testing machine.
  • the strength, elongation and r90 value data can be obtained.
  • the value range of r90 value adopts the tensile strain range of 4-6%;
  • the hardness test uses a microhardness tester to test the HV hardness, with a loading load of 150g and a loading time of 15s.
  • Table 4 lists the relevant performance parameters of the high r-value high strength steel plates of Examples 1-6 and the comparative steel plates of Comparative Examples 1-2.
  • the upper and lower surface hardness in the table have two columns of data, representing the hardness values of the upper surface and the lower surface respectively.
  • the tensile strength of the high-strength steel plates of Examples 1-6 of the present invention is ⁇ 985 MPa
  • the r value r 90 thereof is ⁇ 1.2
  • the microhardness of the upper and lower surface layers is HV ⁇ 130
  • the microhardness of the middle layer is HV ⁇ 330
  • Comparative Example 1 The amount of Ti added in Comparative Example 1 does not satisfy the requirements of the present invention, and even if the process of the present invention is used, a high r value cannot be achieved.
  • Comparative Example 2 uses the ingredients of Example 2 of the present invention, it does not use the annealing process of the present invention, but uses a conventional quenching and tempering process, so a high r value cannot be obtained.

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)
  • Chemical Kinetics & Catalysis (AREA)
  • General Chemical & Material Sciences (AREA)

Abstract

本发明公开了一种高r值超高强钢板，其在厚度方向上包括上表层、中间层和下表层；其中上表层和下表层的微观组织的主体是铁素体，上表层和下表层的含碳量均≤0.025％；上表层和下表层中的化学元素质量百分含量分别满足：Ti-3.42N-3.98C≥0；中间层的微观组织主体为回火马氏体和贝氏体的至少其中之一。相应地，本发明还公开了一种高r值超高强钢板的制造方法。本发明的钢板在具有超高强度的同时还具有高r值，该高r值超高强钢板在变形过程中的厚度减薄更为轻微，从而不论是全局性拉伸变形还是局部性变形中，都变现出较好的成形性。

Description

一种高r值的超高强钢板及其制造方法 技术领域
本发明涉及一种钢板及其制造方法，尤其涉及一种高强钢板及其制造方法。
背景技术
r值代表钢板在拉延成形的变形过程中，板厚抗减薄能力的高低。在抗拉强度350MPa以下的高成形性深冲钢板(如无间隙原子钢)中，r值是重要的技术指标之一，r值高意味着钢板具有更优的拉延成形性能。
然而现有的高强钢钢板，由于不用于深冲，并未关注钢板的r值及其所表征的性能。
例如，公开号为CN101768695A，公开日为2010年7月7日，名称为“1000MPa级Ti微合金化超细晶冷轧双相钢及其制备工艺”的中国专利文献公开了，其化学成分为：0.03-0.2％C，0.2-0.8％Si，1.2-2.0％的Mn，Ti：0.03-0.15％，≤0.02％P，S≤0.015，0.02-0.15％Al，其它为Fe和不可避免杂质组成。经热轧、冷轧后，在临界区退火，冷却速率小于50℃/s的情况下，得到980MPa以上的冷轧双相钢。可以看出，该高强钢并未关注钢板的r值。
例如，公开号为CN101363099A，公开日为2009年2月11日，名称为“一种抗拉强度1000MPa级冷轧双相钢板及制备方法”的中国专利文献公开了，C：0.14-0.21％，Si：0.4-0.9％，Mn：1.5-2.1％，P：≤0.02％，S≤0.01％，Nb：0.001-0.05％，V：0.001-0.02％，经热轧冷轧后，在760-820℃间保温，冷速40-50℃/s，在240-320℃过时效180-300s。该高强钢并未关注钢板的r值。
熊自柳等人在《高强度双相钢的成形性能》(金属热处理2021年第46卷第5期)中介绍了一种抗拉强度为980MPa的双相钢，但是其r值低于0.9。
葛德龙等人在《基于三点弯曲的1000MPa级超高强马氏体钢最大弯曲角度》(塑性工程学报，2014年第21卷第4期)中介绍了1000-1400MPa的马氏体钢，其r值低于0.7。
发明内容
本发明的目的之一在于提供一种高r值超高强钢板，这种钢板在具有超高强度的同时，还具有高r值，该高r值超高强钢板在变形过程中的厚度减薄更为轻微，从而不论是全局性拉伸变形还是局部性变形中，都表现出较好的成形性。
在本发明所述的高r值超高强钢板中，r值代表钢板在拉延成形的变形过程中，板厚抗减薄能力的高低。高r值的钢板，在拉伸变形时，钢板的厚度不易减薄，因此不易发生断裂。对于超高强钢板来说，其在局部发生剧烈变形时，如果r值高，局部不易发生厚度减薄，可以减缓缩颈过程，因此有利于钢板的成形性。
为了实现上述目的，本发明提出了一种高r值超高强钢板，在厚度方向上其包括上表层、中间层和下表层；其中上表层和下表层的微观组织的主体是铁素体，上表层和下表层的含碳量均≤0.025％，所述上表层和下表层中的化学元素质量百分含量均分别满足：Ti-3.42N-3.98C≥0；中间层的微观组织包括铁素体和回火马氏体。
本发明的核心之一就在于，通过使得上、下表层中的化学元素质量百分含量均分别满足Ti-3.42N-3.98C≥0，上表层和下表层的含碳量均≤0.025％，从而实现高r值。同时钢板整体的高强度则主要靠中间层的组织来保证。在一些实施方案中，上表层和下表层的含碳量均为0.02-0.025％。所述上表层和下表层中的化学元素质量百分含量均分别满足：Ti-3.42N-3.98C≥0.060，优选≥0.068。在一些实施方案中，所述上表层和下表层中的化学元素质量百分含量均分别满足：0.16≥Ti-3.42N-3.98C≥0.06。
进一步地，在本发明所述的高r值超高强钢板中，所述上表层和下表层中的铁素体的平均晶粒直径≤15微米。在一些实施方案中，在本发明所述的高r值超高强钢板中，所述上表层和下表层中的铁素体的平均晶粒直径可≤10微米。在一些实施方案中，在本发明所述的高r值超高强钢板中，所述上 表层和下表层中的铁素体的平均晶粒直径为6-10微米。
进一步地，在本发明所述的高r值超高强钢板中，所述上表层和下表层中的铁素体的体积相比例各自≥97％。在一些实施方案中，所述上表层和下表层中的铁素体的体积相比例为97-99％。
进一步地，在本发明所述的高r值超高强钢板中，所述中间层的微观组织还可以包括贝氏体。
更进一步地，在本发明所述的高r值超高强钢板中，无论中间层是否含有贝氏体，回火马氏体+贝氏体的体积相比例≥40％。在一些实施方案中，中间层含有铁素体、以及回火马氏体和贝氏体中的至少一种。在一些实施方案中，中间层中铁素体的体积相比例为55-60％，优选为57-60％。在一些实施方案中，中间层中回火马氏体的体积相比例为40-45％，优选为40-43％。在一些实施方案中，中间层中贝氏体的体积相比例为0-5％，优选为0-3％。
进一步地，在本发明所述的高r值超高强钢板中，所述上表层和下表层的铁素体中含有碳化物，所述碳化物的类型为Ti(C，N)和Nb(C，N)。
进一步地，在本发明所述的高r值超高强钢板中，其上表层和下表层的厚度均分别为100-200微米。
进一步地，在本发明所述的高r值超高强钢板中，其上表层和下表层的厚度之和不超过钢板总厚度的40％。
进一步地，本发明所述的高r值超高强钢板的厚度为0.8-2.5mm，优选为1.2-2.0mm。
进一步地，本发明所述的高r值超高强钢板，其含有Fe和不可避免的杂质元素，此外还含有质量百分含量如下的下述各化学元素：
C：0.08-0.20％；Si：0.01-1.7％；Mn：0.7-2.7％；Ti：0.16-0.26％；Nb：0.015-0.05％；Al：0.02-0.06％；N≤0.005％。
进一步地，在本发明所述的高r值超高强钢板，其各化学元素质量百分比为：
C：0.08-0.20％；Si：0.01-1.7％；Mn：0.7-2.7％；Ti：0.16-0.26％；Nb：0.015-0.05％；Al：0.02-0.06％；N≤0.005％；余量为Fe和不可避免的杂质。
在一些实施方式中，本发明所述高r值超高强钢板的各化学元素的设计原理具体如下所述：
C：在本发明所述的高r值高强钢板中，C可以通过影响马氏体硬度来提高钢的强度。若钢中含碳量过低，则淬火后钢的强度较低；而钢中含碳量越高，淬火后钢的强度越高。但过高的碳不利于获得高r值。因此，在本发明所述的高r值超高强钢板中将C元素含量控制在0.08-0.20％之间。
Si：本发明所述的高r值高强钢板中，Si可以起到固溶强化、抑制Fe3C析出、促进残余奥氏体形成等作用。基于此，本发明控制Si含量为0.01-1.7％。在一些实施方案中，Si的含量为0.4-1.7％。
Mn：Mn是提高钢的淬硬性的主要元素，Mn的含量与采用的冷却方式的冷却能力相匹配。当Mn的含量过低时，钢板无法淬硬，从而无法得到较高强度，而若Mn含量过高，则对碳当量不利。因此，在本发明所述的高r值超高强钢板中将Mn元素含量控制在0.7-2.7％之间。在一些实施方案中，Mn的含量为1.0-2.7％
Ti：本发明所述的高r值高强钢板中，Ti的含量和本发明所述的r值关系密切，Ti含量过低，则不足以固定上下表层的全部C、N，导致r值偏低；Ti含量过高，则Ti过剩。因此，在本发明所述的高r值超高强钢板中将Ti元素控制在0.16-0.26％之间。
Nb：本发明所述的高r值高强钢板中，Nb起到辅助固定C、N的作用，形成Nb(C，N)，确保上下表层区域中的C、N都不处于固溶状态。因此，在本发明所述的高r值超高强钢板中将Nb元素控制在0.015-0.05％之间。
Al：在本发明中的主要作用是冶炼过程中脱氧，其含量范围可以为：0.02-0.06％。
N：在本发明中，N不利于在上下表层形成Ti-3.42N-3.98C≥0的成分比例，不利于在局部形成无间隙原子的状态，因此需要控制N≤0.005％。
在本发明中，不可避免的杂质主要是S和P，在工艺条件允许的情况下，期望其含量越低越好。在本发明所述的高r值超高强钢板中，P元素的含量≤0.015％。在本发明所述的高r值超高强钢板中，S元素的含量≤0.003％。
可以这样理解，在本发明所述的高r值高强钢板中，基本成分为：C：0.08-0.20％；Si：0.01-1.7％；Mn：0.7-2.7％。该基本成分构成钢板强度的基础，但基于这些基本成分并不能获得高r值。为了获得高r值，本发明所述的钢板中还含有：Ti：0.16-0.26％，Nb：0.015-0.05％，N≤0.005％。
进一步地，在本发明所述的高r值超高强钢板中，其化学元素还含有Cr、Mo、B的至少其中一种；其中，B≤0.005％，Cr+Mo≤0.5％。在一些实施方案中，含有时，B的含量不超过0.005％，Cr的含量不超过0.3％，Mo的含量不超过0.2％。
在本发明所述的适量的高r值超高强钢板中，Cr、Mo、B合金元素可以有效配合锰，从而起到提高淬透性的效果，以使钢的淬透性和快冷冷却能力相匹配，以进一步避免无法淬硬或者碳当量过高。
进一步地，本发明所述的高r值超高强钢板的抗拉强度≥980MPa，其r值r90≥1.2，上表层和下表层的显微硬度HV≤130，中间层的显微硬度HV≥330。优选地，本发明所述的高r值超高强钢板的抗拉强度≥1020MPa。在一些实施方案中，本发明所述的高r值超高强钢板的抗拉强度可为980-1100MPa，优选为980-1060MPa。
在一些实施方案中，本发明所述的高r值超高强钢板的r值r90为1.2-1.4。
在一些实施方案中，本发明所述的高r值超高强钢板的上表层和下表层的显微硬度HV各自独立为100-130。
在一些实施方案中，本发明所述的高r值超高强钢板的中间层的显微硬度HV为330-380。
在一些实施方案中，本发明所述的高r值超高强钢板的延伸率≥12％，如12-16％或13-15％。
在一些实施方案中，本发明所述的高r值超高强钢板的屈服强度≥570MPa，如570-660MPa。
相应地，本发明的另一目的在于提供一种高r值超高强钢板的制造方法，该方法通过优化工艺，可以制得高r值超高强钢板。
为了实现上述目的，本发明还提出了如上文所述的高r值超高强钢板的制造方法，包括步骤：冶炼和铸造；热轧；酸洗后冷轧；退火；其中所述退火步骤包括：
高温增湿脱碳过程：在退火炉内喷水汽增湿，同时控制钢板加热温度为880-950℃保持100-300s，控制退火炉内露点为0℃以上；
普通缓冷过程：钢板从高温增湿的温度，以3-10℃/s的冷速冷却到730-770℃之间；
准静态缓冷过程：钢板从730-770℃缓冷到680-720℃，冷却温度区间控制在40-60℃，冷却速度0.03-0.1℃/s，保证缓冷过程持续500s以上；
快冷过程：钢板以≥50℃/s的冷速快速冷却到300℃以下；
过时效过程：钢板在240-320℃回火150-500s。
本发明所述的制造方法所采用的工艺，尤其是独特的退火工艺，对于本发明所述的超高强度钢板获得高r值是至关重要的。
其中，脱碳的方法为炉内喷水汽增湿，在高温下发生如下反应，促进钢板脱碳反应：C+H2O＝CO+H2；C+2H2O＝CO2+2H2。钢板加热温度为880～950℃之间，并保持100-300s，根据单位时间通过带钢的重量和需要脱碳层的厚度，控制增湿的水量，保证退火炉内露点为0℃以上。在一些实施方案中，退火炉内露点为0-20℃。该高温增湿脱碳工艺使得钢板的上表层和下表层分别成为厚度为100-200微米的初步脱碳层(即上表层和下表层均为脱碳层)。在这个过程中，如果钢板加热温度过低，会导致脱碳不足，钢板加热温度过高，则会晶粒过分长大，因此，在高温增湿脱碳过程中控制钢板加热温度为880-950℃，有以利于钢板的脱碳反应。而选择炉内露点在0℃以上的原因是因为炉内露点低，则钢板的脱碳反应慢，不足以得到足够厚度的脱碳层。
在上述高温增湿脱碳过程之后，钢板上下表层的含碳量明显下降，可以下降到≤0.05％的水平。此时钢板上下表层的含碳量仍然过高，奥氏体直接冷却转变成铁素体，仍会析出Fe3C。为了进一步降低钢板上下表层的含碳量，在铁碳相图的Ac1温度以下进一步进行准静态缓冷过程处理。在该准静态缓冷过程中，钢板以0.03-0.1℃/s的速度从730-770℃缓冷到680-720℃，整个缓冷过程温度区间控制在40-60℃，持续时间≥500s。在一些实施方案中，缓冷持续时间为500-1500s。在该温度区间内，经过初步脱碳的钢板上下表层率先发生奥氏体分解，形成铁素体，在缓冷过程中，铁素体中的碳向未发生转变的高碳区(中间层)的奥氏体中扩散，从而进一步促使上、下表层的含碳量下降，在缓冷时间充分的情况下，上、下表层的含碳量可降低到接近工业纯铁的含碳量水平，从而使发明钢的上下表层区域满足Ti-3.42N-3.98C≥0。
在本发明所述的制造方法中，从730-770℃缓冷到680-720℃至关重要，脱碳后表层的奥氏体退碳含量降低，Ac3上升，从而在从730-770℃缓冷到680-720℃的过程中率先发生奥氏体→铁素体转变。而本发明钢板的上、下表 层局部含碳量降低到≤0.025％时，钢板中的Ti、C、N含量就能满足Ti-3.42N-3.98C≥0的水平，从而实质上在上、下表层形成了无间隙原子的结构，大大提高上、下表层的r值，进而提高了钢板整体的r值。
缓冷结束后，钢板以≥50℃/s的冷速快速冷却到300℃以下，并在240-320℃回火150-500s。在一些实施方案中，快冷的冷速为50-150℃/s。在一些实施方案中，快冷的终止温度为240-300℃。该冷却过程中，中间层的奥氏体转变成马氏体组织，在一些实施方式中也可以含有少量贝氏体和残余奥氏体组织，从而实现超高强度，例如≥980MPa的抗拉强度。
在一些实施方案中，过时效温度为240-300℃。
本发明所述的制造方法中的冶炼、铸造、热轧和冷轧工艺均可以采用常规工艺，本发明对这些工艺并未进行特别地改进。示例性的热轧步骤中，钢板加热出炉温度为1220-1280℃，终轧温度为870-930℃，卷取温度为570-630℃。示例性的冷轧工序中，控制冷轧压下率为40-65％。
本发明所述的高r值高强度钢板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：
本发明通过控制中间层的组织形态来保证超高强度，同时通过控制上、下表层中的化学元素质量百分含量均分别满足Ti-3.42N-3.98C≥0，从而实现高r值，从而实现了超高强度钢板具有高r值。
从成分设计的角度来说，本发明通过添加适量的Ti并辅助以Nb固定N、C，可以实质性地在上下表层实现无间隙原子的固溶体状态，从而大大提高了表层的r值，进而影响到钢板整体的r值，使980MPa的冷轧高强钢钢板的r值达到1.2以上，从而在变形过程中减薄率大大减轻，从而获得更优的成形性。
从制造工艺的角度来说，本发明所述的高r值高强度钢板通过采用独特的退火工艺，经过初步高温增湿脱碳工艺和长时间缓冷奥氏体→铁素体转变，使得碳含量再分配，从而能够使超高强钢板的上、下表层的含碳量降低到0.025％以下。
附图说明
图1示意性地显示了本发明所述的高r值高强度钢板的制造方法在一种 实施方式下的退火工艺曲线。
具体实施方式
下面将结合说明书附图和具体的实施例对本发明所述的高r值高强度钢板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1-6的高r值高强度钢板和对比例1-2的对比钢板通过以下步骤制得：
(1)冶炼：铁水经过转炉进行冶炼，并经过精炼过程进一步优化成分；
(2)连铸获得板坯，各实施例和对比例的板坯成分含量如表1所示；
(3)热轧：采用常规的热轧工艺，具体而言，控制钢板加热出炉温度为1220-1280℃，终轧温度为870-930℃，卷取温度为570-630℃；
(4)酸洗后冷轧：控制冷轧压下率为40-65％；
(5)退火，具体包括：
高温增湿脱碳过程：在退火炉内喷水汽增湿，同时控制钢板加热温度为880-950℃保持100-300s，控制退火炉内露点为0℃以上；
普通缓冷过程：钢板从高温增湿的温度，以3-10℃/s的冷速冷却到730-770℃之间；
准静态缓冷过程：钢板从730-770℃缓冷到680-720℃，冷却温度区间控制在40-60℃，冷却速度0.03-0.1℃/s，保证缓冷过程持续500s以上；
快冷过程：钢板以≥50℃/s的冷速快速冷却到300℃以下；
过时效过程：钢板在240-320℃回火150-500s。
其中，实施例1-6以及对比例1的退火工艺的参数列于表2-1中。对比例2未采用本发明的退火工艺，而是采用了常规的淬火+回火工艺，具体工艺参数列于表2-2中。
表1列出了本发明实施例1-6和对比例1-2的化学元素的质量百分配比。
表1.(wt％，余量为Fe和除了S、P、N以外其他不可避免的杂质)

表2-1列出了实施例1-6的高r值高强度钢板和对比例1-2的对比钢板的退火步骤的具体工艺参数。
表2-1
表2-2
对实施例1-6以及对比例1-2的钢板进行取样，采用GDS(辉光放电光谱仪)对钢板上下表层的含碳量及其它组分含量进行检测，并用金相显微镜对其微观组织进行观察和检测，并将观察和检测结果列于表3中。
表3.
注：当上、下表层中的Ti-3.42N-3.98C值大于0时，说明上下表层无间隙原子，而当Ti-3.42N-3.98C值小于0时，说明上下表层有间隙原子。表中的两列数据分别表示上表层和下表层的数值。
为了对本发明的实施效果进行验证，对实施例1-6以及对比例1-2的钢 板进行取样，并检测其性能，并将检测结果列于表4中。其中：
力学性能测试沿钢板的垂直于轧向的方向进行拉伸测试。按照标距长度50mm，标距区域宽度25mm的JIS5#标准板状拉伸试样，在拉伸试验机上进行检测，检测过程中可以获得强度、延伸率和r90值的数据，r90值的取值范围采用4-6％拉伸应变范围；
硬度检测采用显微硬度仪测试HV硬度，加载载荷150g，加载时间15s。
表4列出了实施例1-6的高r值高强钢板以及对比例1-2的对比钢板的相关性能参数。
表4.
注：表中上、下表层硬度具有两列数据分别表示上表层和下表层的硬度数值。
从表4可以看出，本发明实施例1-6的高强钢板抗拉强度均≥985MPa，其r值r90≥1.2，上表层和下表层的显微硬度HV≤130，中间层的显微硬度HV≥330，这说明了本发明所述的高强钢板可以兼具高强度和高成形性。
对比例1的Ti的添加量不满足本发明，即使其采用了本发明的工艺，也无法做到高r值。
对比例2虽然采用了本发明实施例2的成分，但是其未采用本发明的退火工艺，而是采用了常规的淬火回火工艺，因此也无法获得高r值。
需要说明的是，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. 一种高r值超高强钢板，其特征在于，在厚度方向上包括上表层、中间层和下表层；其中上表层和下表层的微观组织的主体是铁素体，上表层和下表层的含碳量均≤0.025％；所述上表层和下表层中的化学元素质量百分含量分别满足：Ti-3.42N-3.98C≥0；中间层的微观组织包括铁素体和回火马氏体。
  2. 如权利要求1所述的高r值超高强钢板，其特征在于，所述上表层和下表层中的铁素体的平均晶粒直径≤15微米。
  3. 如权利要求1所述的高r值超高强钢板，其特征在于，所述上表层和下表层中铁素体的体积相比例≥97％。
  4. 如权利要求1所述的高r值超高强钢板，其特征在于，所述中间层的微观组织还包括贝氏体。
  5. 如权利要求4所述的高r值超高强钢板，其特征在于，所述中间层的回火马氏体+贝氏体的体积相比例≥40％；优选地，所述中间层中，铁素体的体积相比例为55-60％，回火马氏体的体积相比例为40-45％，贝氏体的体积相比例为0-5％。
  6. 如权利要求1所述的高r值超高强钢板，其特征在于，所述上表层和下表层的铁素体中含有碳化物，所述碳化物的类型为Ti(C，N)和Nb(C，N)。
  7. 如权利要求1所述的高r值超高强钢板，其特征在于，其上表层和下表层的厚度均分别为100-200微米。
  8. 如权利要求1所述的高r值超高强钢板，其特征在于，其厚度为0.8-2.5mm。
  9. 如权利要求1所述的高r值超高强钢板，其特征在于，其上表层和下表层的厚度之和不超过钢板总厚度的40％。
  10. 如权利要求1所述的高r值超高强钢板，其含有Fe和不可避免的杂质元素，其特征在于，其还含有质量百分含量如下的下述各化学元素：
    C：0.08-0.20％；Si：0.01-1.7％；Mn：0.7-2.7％；Ti：0.16-0.26％；Nb：
    0.015-0.05％；Al：0.02-0.06％；N≤0.005％。
  11. 如权利要求10所述的高r值超高强钢板，其特征在于，其各化学元素质 量百分比为：
    C：0.08-0.20％；Si：0.01-1.7％；Mn：0.7-2.7％；Ti：0.16-0.26％；Nb：
    0.015-0.05％；Al：0.02-0.06％；N≤0.005％；余量为Fe和不可避免的杂质。
  12. 如权利要求10或11所述的高r值超高强钢板，其特征在于，其化学元素还含有Cr、Mo、B的至少其中一种；其中，B≤0.005％，Cr+Mo≤0.5％；优选地，B≤0.005％，Cr≤0.3％，Mo≤0.2％。
  13. 如权利要求1所述的高r值超高强钢板，其特征在于，其抗拉强度≥980MPa，其r值r90≥1.2，上表层和下表层的显微硬度HV≤130，中间层的显微硬度HV≥330。
  14. 一种如权利要求1-13中任意一项所述的高r值超高强钢板的制造方法，其特征在于，包括步骤：冶炼和铸造；热轧；酸洗后冷轧；退火；其中所述退火步骤包括：
    高温增湿脱碳过程：在退火炉内喷水汽增湿，同时控制钢板加热温度为880-950℃保持100-300s，控制退火炉内露点为0℃以上；
    普通缓冷过程：钢板从高温增湿的温度，以3-10℃/s的冷速冷却到730-770℃之间；
    准静态缓冷过程：钢板从730-770℃缓冷到680-720℃，冷却温度区间控制在40-60℃，冷却速度0.03-0.1℃/s，保证缓冷过程持续500s以上；
    快冷过程：钢板以≥50℃/s的冷速快速冷却到300℃以下；
    过时效过程：钢板在240-320℃回火150-500s。
  15. 如权利要求14所述的制造方法，其特征在于，所述方法具有以下一项或多项特征：
    (1)所述热轧为控制钢板加热出炉温度为1220-1280℃，终轧温度为870-930℃，卷取温度为570-630℃；
    (2)所述冷轧的压下率为40-65％；
    (3)所述退火炉内露点为0-20℃；
    (4)所述缓冷持续500-1500s；
    (5)快冷的冷速为50-150℃/s；和
    (6)回火温度为240-300℃。
PCT/CN2024/102541 2023-06-30 2024-06-28 一种高r值的超高强钢板及其制造方法 Ceased WO2025002398A1 (zh)

Priority Applications (2)

Application Number Priority Date Filing Date Title
KR1020267002712A KR20260032574A (ko) 2023-06-30 2024-06-28 r값이 높은 초고강도 강판 및 이의 제조 방법
EP24831025.2A EP4737603A1 (en) 2023-06-30 2024-06-28 Ultrahigh-strength steel plate having high r value and manufacturing method therefor

Applications Claiming Priority (2)

Application Number Priority Date Filing Date Title
CN202310793720.7A CN119220893B (zh) 2023-06-30 2023-06-30 一种高r值的超高强钢板及其制造方法
CN202310793720.7 2023-06-30

Publications (1)

Publication Number Publication Date
WO2025002398A1 true WO2025002398A1 (zh) 2025-01-02

Family

ID=93937599

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/CN2024/102541 Ceased WO2025002398A1 (zh) 2023-06-30 2024-06-28 一种高r值的超高强钢板及其制造方法

Country Status (4)

Country Link
EP (1) EP4737603A1 (zh)
KR (1) KR20260032574A (zh)
CN (1) CN119220893B (zh)
WO (1) WO2025002398A1 (zh)

Citations (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH05195149A (ja) * 1992-01-21 1993-08-03 Nkk Corp 曲げ加工性及び衝撃特性の優れた超高強度冷延鋼板
US20140212684A1 (en) * 2011-07-29 2014-07-31 Nippon Steel & Sumitomo Metal Corporation High-strength galvanized steel sheet excellent in bendability and manufacturing method thereof
CN110114500A (zh) * 2016-12-23 2019-08-09 Posco公司 冲击特性优异的热压成型用镀覆钢板、热压成型部件及它们的制造方法
CN110199045A (zh) * 2017-04-28 2019-09-03 日本制铁株式会社 高强度钢板及其制造方法
US20190390295A1 (en) * 2017-03-30 2019-12-26 Jfe Steel Corporation Hot pressed part and method of manufacturing same
US20200010919A1 (en) * 2017-02-20 2020-01-09 Nippon Steel Corporation High strength steel sheet
CN113195774A (zh) * 2018-12-19 2021-07-30 Posco公司 热成型后的冲击特性优异的热成型用镀覆钢板、热成型部件和它们的制造方法
CN113355590A (zh) * 2020-03-06 2021-09-07 宝山钢铁股份有限公司 一种三层复合组织高强钢板及其制造方法
CN115605625A (zh) * 2020-08-07 2023-01-13 日本制铁株式会社（Jp） 钢板

Family Cites Families (3)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH0756050B2 (ja) * 1989-05-30 1995-06-14 新日本製鐵株式会社 連続焼鈍による非時効・高焼付硬化・プレス加工用高強度冷延鋼板の製造方法
CN103060703B (zh) * 2013-01-22 2015-09-23 宝山钢铁股份有限公司 一种780MPa级冷轧双相带钢及其制造方法
CN116179949B (zh) * 2023-01-19 2025-02-18 鞍钢股份有限公司 汽车用780MPa级超高扩孔性能冷轧复相钢板及其制备方法

Patent Citations (9)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPH05195149A (ja) * 1992-01-21 1993-08-03 Nkk Corp 曲げ加工性及び衝撃特性の優れた超高強度冷延鋼板
US20140212684A1 (en) * 2011-07-29 2014-07-31 Nippon Steel & Sumitomo Metal Corporation High-strength galvanized steel sheet excellent in bendability and manufacturing method thereof
CN110114500A (zh) * 2016-12-23 2019-08-09 Posco公司 冲击特性优异的热压成型用镀覆钢板、热压成型部件及它们的制造方法
US20200010919A1 (en) * 2017-02-20 2020-01-09 Nippon Steel Corporation High strength steel sheet
US20190390295A1 (en) * 2017-03-30 2019-12-26 Jfe Steel Corporation Hot pressed part and method of manufacturing same
CN110199045A (zh) * 2017-04-28 2019-09-03 日本制铁株式会社 高强度钢板及其制造方法
CN113195774A (zh) * 2018-12-19 2021-07-30 Posco公司 热成型后的冲击特性优异的热成型用镀覆钢板、热成型部件和它们的制造方法
CN113355590A (zh) * 2020-03-06 2021-09-07 宝山钢铁股份有限公司 一种三层复合组织高强钢板及其制造方法
CN115605625A (zh) * 2020-08-07 2023-01-13 日本制铁株式会社（Jp） 钢板

Also Published As

Publication number Publication date
CN119220893A (zh) 2024-12-31
EP4737603A1 (en) 2026-05-06
KR20260032574A (ko) 2026-03-09
CN119220893B (zh) 2025-11-14

Similar Documents

Publication Publication Date Title
CN104508163B (zh) 成形性及定形性优异的高强度热浸镀锌钢板及其制造方法
CN104513927B (zh) 一种抗拉强度800MPa级高强度高韧性钢板及其制造方法
US20210207234A1 (en) High-strength steel sheet and method for manufacturing the same
CN106068333A (zh) 高强度低比重钢板及其制造方法
CN111996467A (zh) 一种980MPa级镀锌高强钢及其制备方法
CN105274432A (zh) 600MPa级高屈强比高塑性冷轧钢板及其制造方法
CN108707823A (zh) 超高强度钢板及其制备方法和超高强度钢板制品
US20130160907A1 (en) High strength cold rolled steel sheet having excellent stretch flangeability and method for manufacturing the same
CN110073024A (zh) 弯曲加工性优异的超高强度钢板及其制造方法
CN113584393A (zh) 一种抗拉强度780MPa级双相钢及其生产方法
JPH1060593A (ja) 強度−伸びフランジ性バランスにすぐれる高強度冷延鋼板及びその製造方法
CN108531832A (zh) 一种800MPa级高屈强比冷轧钢板及其制造方法
JP5842748B2 (ja) 冷延鋼板およびその製造方法
CN113025882B (zh) 一种热基镀锌铁素体贝氏体高强钢板及其制备方法
CA2546003A1 (en) High-rigidity high-strength thin steel sheet and method for producing same
US20250354244A1 (en) 800mpa-grade high-hole-expansion hot-dip galvanized steel plate and manufacturing method therefor
CN117305694B (zh) 一种800MPa级热基镀锌复相钢板及其制造方法
EP4353862A1 (en) Hot-dip galvanized steel plate and manufacturing method therefor
CN107406939B (zh) 高强度冷轧钢板及其制造方法
CN116716550A (zh) 汽车用淬火配分钢及其梯度配分制备方法
CN116043121A (zh) 一种成型性能优异的800MPa级冷轧复相钢及其制备方法
CN102534373B (zh) 一种适于辊压成形的超高强度冷轧钢带及其制造方法
CN117062934A (zh) 钢板、部件以及它们的制造方法
JP2023534825A (ja) 成形性及び加工硬化率に優れた鋼板
CN114341390B (zh) 均匀伸长率和加工硬化率优异的钢板及其制造方法

Legal Events

Date Code Title Description
121 Ep: the epo has been informed by wipo that ep was designated in this application

Ref document number: 24831025

Country of ref document: EP

Kind code of ref document: A1

ENP Entry into the national phase

Ref document number: 1020267002712

Country of ref document: KR

Free format text: ST27 STATUS EVENT CODE: A-0-1-A10-A15-NAP-PA0105 (AS PROVIDED BY THE NATIONAL OFFICE)

WWE Wipo information: entry into national phase

Ref document number: 2024831025

Country of ref document: EP

NENP Non-entry into the national phase

Ref country code: DE

ENP Entry into the national phase

Ref document number: 2024831025

Country of ref document: EP

Effective date: 20260130

ENP Entry into the national phase

Ref document number: 2024831025

Country of ref document: EP

Effective date: 20260130