WO2024111525A1 - Tôle en acier laminée à chaud hautement résistante, et procédé de fabrication de celle-ci - Google Patents

Tôle en acier laminée à chaud hautement résistante, et procédé de fabrication de celle-ci Download PDF

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Publication number
WO2024111525A1
WO2024111525A1 PCT/JP2023/041535 JP2023041535W WO2024111525A1 WO 2024111525 A1 WO2024111525 A1 WO 2024111525A1 JP 2023041535 W JP2023041535 W JP 2023041535W WO 2024111525 A1 WO2024111525 A1 WO 2024111525A1
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amount
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steel sheet
heating
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PCT/JP2023/041535
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English (en)
Japanese (ja)
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寛 長谷川
広志 松田
隼佑 飛田
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Ｊｆｅスチール株式会社
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Publication of WO2024111525A1 publication Critical patent/WO2024111525A1/fr

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    • B—PERFORMING OPERATIONS; TRANSPORTING
    • B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
    • B21B—ROLLING OF METAL
    • B21B1/00—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
    • B21B1/22—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length
    • B21B1/24—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process
    • B21B1/26—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process by hot-rolling, e.g. Steckel hot mill
    • 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
    • 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
    • C22C38/14—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
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur

Definitions

  • the present invention relates to high-strength hot-rolled steel sheets and their manufacturing methods, and in particular to high-strength hot-rolled steel sheets suitable as materials for automotive parts and their manufacturing methods.
  • Patent Document 1 discloses a technique for improving stretch flangeability by setting the processing temperature (post-heating temperature) at 400 to 1000°C.
  • Patent Document 2 discloses a technique for a high-strength hot-rolled steel sheet with a TS of 730 MPa or more.
  • the hot-rolled steel sheet disclosed in Patent Document 2 has a structure in which bainite is the main phase and 80% or more of the total Ti amount is solid-solute Ti. This provides heat treatment hardening with an increase in YS (yield strength) and TS of 100 MPa or more after heat treatment in which the steel sheet is heated to a temperature range of 500°C to the Ac1 transformation point and held for 60 min.
  • Patent Document 3 discloses a technique for a hot-rolled steel sheet with excellent delayed fracture resistance and a TS of 120 kgf/mm2 or more .
  • Patent Document 1 does not take into account the performance of the part, such as strength and toughness after post-heating, and there is room for improvement.
  • the steel structure changes significantly, and when the steel plate (original plate) before post-heating has a high strength of 1180 MPa or more, the impact on the strength becomes significant, so material design that takes into account the impact of post-heating on strength is necessary.
  • the steel plate disclosed in Patent Document 2 does not experience a decrease in strength after heat treatment, but on the contrary, it leads to an excessive increase in strength and there is a problem with the toughness after heat treatment due to the significant fine precipitation of carbides, and there is room for improvement.
  • Patent Document 3 aims to improve the workability of the original plate and improve the delayed fracture resistance without post-heating.
  • the delayed fracture test was evaluated by deep drawing, and no consideration was given to the more severe end surface, or delayed fracture resistance from the end surface processed after post-heating, leaving room for improvement.
  • the present invention was made in consideration of the above circumstances, and aims to provide a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
  • the inventors focused on the precipitation behavior of Ti and Nb after post-heating of hot-rolled steel sheets, and came up with the idea of improving the properties of the steel sheets after heating (after post-heating) by controlling the initial coarse Ti-containing precipitates and Nb-containing precipitates and the amount of dissolved Ti and Nb before post-heating. Furthermore, they focused on the crystal orientation, and came up with the idea of suppressing delayed fracture after post-heating of the punched end surface by forming a structure in which the surface layer region is concentrated in a specific orientation.
  • high strength means that the tensile strength (TS) is 1180 MPa or more and less than 1600 MPa.
  • excellent strength after post-heating means that the decrease in strength of the hot-rolled steel sheet after post-heating is 50 or less in Vickers hardness compared to the strength of the hot-rolled steel sheet (original sheet) before post-heating.
  • being excellent in toughness after post-heating means that in a Charpy impact test using a test piece taken from the hot-rolled steel sheet after post-heating, the ductile fracture surface ratio at ⁇ 20° C. is 50% or more.
  • the thickness of the test piece is 0.6 to 3.0 mm, and when the thickness of the hot-rolled steel sheet exceeds 3.0 mm, the test piece taken from the hot-rolled steel sheet is ground on both sides to a thickness of 3.0 mm, and then subjected to the Charpy impact test.
  • excellent resistance to delayed fracture after post-heating means that no cracks are generated when a rectangular test piece having a sheared end surface, which is taken from the steel sheet after post-heating, is V-bent at 90°, the part opened by springback is tightened with a bolt or the like, and the piece is immersed in hydrochloric acid of pH 3 for 96 hours.
  • post-heating means a heat treatment in which the hot-rolled steel sheet (original sheet) is heated to 400° C. or higher.
  • the present invention has the following configuration.
  • C 0.06 to 0.23%, Si: 0.1 to 3.0%, Mn: 1.5 to 3.5%, P: more than 0% and not more than 0.050%; S: more than 0% and 0.0050% or less; Al: more than 0% and not more than 1.5%; N: more than 0% and not more than 0.010%; O: more than 0% and 0.003% or less; Contains Ti and Nb in total 0.040 to 0.200%; The balance is Fe and unavoidable impurities,
  • the steel structure has martensite and/or lower bainite as a main phase, and the volume fraction of retained austenite is less than 3%; the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb, that is, (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb), is 0.300 or more and less than 0.800;
  • the total amount of Ti and Nb present as precipitates having
  • the composition further comprises, in mass%, Cr: 0.005 to 2.0%, Ni: 0.005 to 2.0%, Mo: 0.005 to 1.0%, V: 0.005 to 0.5%, B: 0.0002 to 0.0050%, Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050% Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%
  • the high strength hot rolled steel sheet according to [1] comprising one or more selected from the following: [3] A method for producing a high strength hot rolled steel sheet according to the above [1] or [2], A slab having the above-mentioned composition is heated to a temperature range of 1150 to 1300°C and held at that temperature range for 0.2 to 3.5 hours; Next, when hot rolling is performed, The total reduction in the temperature range of 1080°C or more is 80 to 90%, the total reduction in the temperature range of 900°C or less is 20% or more, and the reduction per pass at T
  • the steel is allowed to cool for 1.0 s or more.
  • the temperature is cooled at an average cooling rate of 50°C/s or more up to 550°C, and the time from reaching 550°C to starting quenching is set to 0.5 to 4.0 s.
  • the steel sheet is quenched to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more, and then coiled at the coiling temperature.
  • T(°C) 800+1000[Ti]+2500[Nb]
  • [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
  • the present invention provides a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
  • the present invention it is possible to obtain a high-strength hot-rolled steel sheet which is suitable as a material for automobile parts and which has excellent strength, toughness and delayed fracture resistance after post-heating or after post-heat processing.
  • a high-strength hot-rolled steel sheet of the present invention it is possible to obtain products such as high-strength automobile parts that exhibit high strength, good toughness, and excellent delayed fracture resistance even after heat treatment is performed to improve workability and fatigue properties.
  • the high-strength hot-rolled steel sheet of the present invention may be either a black skin as hot-rolled, or a hot-rolled steel sheet called a white skin which is further pickled after hot rolling.
  • the high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 0.6 mm or more.
  • the high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 10.0 mm or less. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 1.0 mm or more.
  • the thickness is more preferably 6.0 mm or less.
  • the width of the high-strength hot-rolled steel sheet of the present invention is preferably 500 mm or more, more preferably 700 mm or more.
  • the width of the high-strength hot-rolled steel sheet of the present invention is preferably 1800 mm or less, more preferably 1400 mm or less.
  • the high-strength hot-rolled steel sheet of the present invention has a specific chemical composition and a specific steel structure.
  • the chemical composition and steel structure will be explained in that order.
  • the composition of the high-strength hot-rolled steel sheet of the present invention is, in mass%, C: 0.06-0.23%, Si: 0.1-3.0%, Mn: 1.5-3.5%, P: 0.050% or less (excluding 0%), S: 0.0050% or less (excluding 0%), Al: 1.5% or less (excluding 0%), N: 0.010% or less (excluding 0%), O: 0.003% or less (excluding 0%), Ti and Nb in total 0.040-0.200%, and the balance consisting of Fe and unavoidable impurities.
  • C 0.06 to 0.23%
  • C is an element that is effective in increasing TS by generating and strengthening martensite and lower bainite, and in suppressing strength reduction after post-heating by combining with Ti, Nb, N, etc. to generate precipitates. If the C content is less than 0.06%, such effects cannot be sufficiently obtained, and the TS of the steel plate (original plate) of 1180 MPa or more or excellent strength after post-heating cannot be obtained. On the other hand, if the C content exceeds 0.23%, the decrease in toughness after post-heating becomes significant, and excellent toughness after post-heating cannot be obtained. Therefore, the C content is set to 0.06 to 0.23%.
  • the C content is preferably set to 0.07% or more.
  • the C content is preferably set to 0.22% or less, and more preferably set to 0.20% or less.
  • Si 0.1 to 3.0% Silicon is an element effective in solution strengthening of steel and suppressing the decrease in strength after post-heating. To obtain such an effect, the silicon content must be 0.1% or more. On the other hand, if the silicon content exceeds 3.0%, polygonal ferrite is excessively formed and the steel structure of the present invention cannot be obtained. Therefore, the silicon content is set to 0.1 to 3.0%.
  • the silicon content is preferably set to 0.2% or more.
  • the silicon content is preferably set to 2.0% or less, more preferably 1.5% or less.
  • Mn 1.5 to 3.5%
  • Mn is an element effective in suppressing ferrite and upper bainite and generating lower bainite and martensite. If the Mn content is less than 1.5%, this effect is not sufficiently obtained, and polygonal ferrite, upper bainite, etc. are generated, and the microstructure of the present invention cannot be obtained. On the other hand, if the Mn content exceeds 3.5%, the deterioration of toughness and delayed fracture resistance becomes significant, and excellent toughness and delayed fracture resistance after post-heating cannot be obtained. Therefore, the Mn content is set to 1.5 to 3.5%.
  • the Mn content is preferably 1.6% or more.
  • the Mn content is preferably 3.0% or less, more preferably 2.5% or less.
  • P more than 0% and 0.050% or less P reduces the toughness and delayed fracture resistance after post-heating, so it is desirable to reduce the amount as much as possible.
  • a P content of up to 0.050% is acceptable. Therefore, the P content is set to 0.050% or less.
  • the P content is preferably set to 0.030% or less. There is no particular lower limit, and the P content may be more than 0%, but if the P content is less than 0.001%, the production efficiency decreases, so the P content is preferably 0.001% or more.
  • the S content can be up to 0.0050%. Therefore, the S content is set to 0.0050% or less.
  • the S content is preferably set to 0.0030% or less, more preferably set to 0.0020% or less, and even more preferably set to 0.0015%.
  • the S content may be more than 0%, but if the S content is less than 0.0002%, the production efficiency decreases, so the S content is preferably 0.0002% or more.
  • Al more than 0% and not more than 1.5% Al acts as a deoxidizer, and is preferably added in the deoxidization process.
  • the Al content may be more than 0%, but from the viewpoint of using it as a deoxidizer, the Al content is preferably 0.01% or more.
  • the Al content is allowed up to 1.5%. Therefore, the Al content is set to 1.5% or less.
  • the Al content is preferably set to 0.50% or less, more preferably 0.20% or less.
  • N more than 0% and 0.010% or less N generates TiN and NbC and inhibits the precipitation of fine TiC, NbC, etc., so it is preferable to reduce the amount as much as possible.
  • an N content of up to 0.010% is acceptable. Therefore, the N content is set to 0.010% or less.
  • the N content is preferably set to 0.007% or less. There is no particular lower limit, and the N content may be more than 0%, but if the N content is less than 0.0005%, the production efficiency decreases, so the N content is preferably 0.0005% or more.
  • O more than 0% and 0.003% or less O reduces toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible.
  • an O content of up to 0.003% is acceptable. Therefore, the O content is set to 0.003% or less.
  • the O content is preferably set to 0.002% or less. There is no particular lower limit, and the O content may be more than 0%, but if the O content is less than 0.0002%, production efficiency decreases, so the O content is preferably 0.0002% or more.
  • Ti and Nb are the most important elements in the present invention, and are necessary elements for obtaining excellent strength, toughness, and delayed fracture resistance properties after post-heating by generating appropriate fine precipitates such as TiC and NbC after post-heating. If the total content of Ti and Nb is less than 0.040%, such effects are not sufficiently obtained, and excellent strength after post-heating is not obtained. On the other hand, if the total content of Ti and Nb exceeds 0.200%, the amount of coarse precipitates containing Ti and Nb increases, which leads to a decrease in delayed fracture resistance properties after post-heating, and the precipitates after post-heating become excessive, making it impossible to obtain excellent toughness after post-heating.
  • the total content of Ti and Nb is set to 0.040 to 0.200%.
  • the total content of Ti and Nb is preferably 0.050% or more, more preferably 0.060% or more.
  • the total content of Ti and Nb is preferably 0.160% or less, more preferably 0.120% or less.
  • the total content of Ti and Nb needs to be within the above range, and the content of either one of them may be 0%.
  • the above components are the basic components of the high-strength hot-rolled steel sheet of the present invention.
  • the high-strength hot-rolled steel sheet of the present invention contains the above components, with the remainder being Fe and unavoidable impurities.
  • the high-strength hot-rolled steel sheet of the present invention can further contain one or more selected from Cr: 0.005-2.0%, Ni: 0.005-2.0%, Mo: 0.005-1.0%, V: 0.005-0.5%, B: 0.0002-0.0050%, Ca: 0.0001-0.0050%, REM: 0.0001-0.0050%, Cu: 0.005-0.5%, Sb: 0.0010-0.10%, Sn: 0.0010-0.10%.
  • Cr 0.005 to 2.0% Cr is an element effective in suppressing ferrite and generating lower bainite and martensite.
  • the Cr content is preferably 0.005% or more.
  • the corrosion resistance may be significantly decreased, so when Cr is contained, the Cr content is preferably 2.0% or less.
  • the Cr content is more preferably 0.1% or more.
  • the Cr content is more preferably 0.8% or less.
  • Ni 0.005 to 2.0%
  • Ni is an element effective in suppressing ferrite and generating lower bainite and martensite.
  • the Ni content is preferably 0.005% or more.
  • the Ni content exceeds 2.0%, a large amount of residual ⁇ is formed, which may lead to a decrease in toughness after post-heating, so when Ni is contained, the Ni content is preferably 2.0% or less.
  • the Ni content is more preferably 0.05% or more.
  • the Ni content is more preferably 0.8% or less, and further preferably 0.5% or less.
  • Mo 0.005 to 1.0%
  • Mo is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
  • the Mo content is preferably 0.005% or more.
  • the Mo content exceeds 1.0%, the generation of Mo-based precipitates becomes significant, which may lead to a decrease in toughness after post-heating, so when Mo is contained, the Mo content is preferably 1.0% or less.
  • the Mo content is more preferably 0.05% or more.
  • the Mo content is more preferably 0.50% or less.
  • V 0.005 to 0.5%
  • V is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
  • the V content is preferably 0.005% or more.
  • the V content exceeds 0.5%, the generation of V-based precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when V is contained, the V content is preferably 0.5% or less.
  • the V content is more preferably 0.01% or more.
  • the V content is more preferably 0.1% or less.
  • B 0.0002 to 0.0050%
  • B is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
  • the B content is preferably 0.0002% or more.
  • the B content exceeds 0.0050%, B-based compounds increase, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when B is contained, the B content is preferably 0.0050% or less.
  • the B content is more preferably 0.0005% or more.
  • the B content is more preferably 0.0040% or less.
  • Ca and REM are each an element effective in improving toughness and delayed fracture resistance after post-heating by controlling the shape of inclusions.
  • the respective contents are preferably 0.0001% or more.
  • the Ca and REM contents are preferably 0.0050% or less.
  • the Ca content is more preferably 0.0005% or more.
  • the Ca content is more preferably 0.0030% or less.
  • the REM content is more preferably 0.0005% or more. Moreover, the REM content is more preferably 0.0030% or less.
  • REM is a collective term for Sc, Y, and 15 other elements ranging from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.
  • Cu 0.005 to 0.5%
  • Sb 0.0010 to 0.10%
  • Sn 0.0010 to 0.10%
  • Cu, Sb, and Sn are each an element that is effective in retarding the corrosion reaction and improving the delayed fracture resistance after post-heating.
  • the Cu content is 0.005% or more
  • the Sb content is 0.0010% or more
  • the Sn content is 0.0010% or more, respectively.
  • the Cu content exceeds 0.5%, the generation of Cu precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when Cu is contained, it is preferable that the Cu content is 0.5% or less.
  • the Sb and Sn contents each exceed 0.10%, the grain boundary embrittlement effect becomes excessive, which may lead to a decrease in delayed fracture resistance, so when Sb and Sn are contained, it is preferable that the Sb and Sn contents each are 0.10% or less.
  • the Cu content is more preferably 0.05% or more. Furthermore, the Cu content is more preferably 0.3% or less.
  • the Sb content is more preferably 0.0050% or more.
  • the Sb content is more preferably 0.050% or less.
  • the Sn content is more preferably 0.0050% or more.
  • the Sn content is more preferably 0.050% or less.
  • the effect of the present invention is not impaired even if the content of Cr, Ni, Mo, V, B, Ca, REM, Cu, Sb, and Sn is less than the lower limit value. Therefore, when the content of these components is less than the lower limit value, these elements are treated as unavoidable impurities.
  • the present invention may further contain one or more of Mg, As, W, Ta, Pb, Zr, Hf, Te, Bi, and Se in a total amount of 0.3% or less by mass. It is preferable to limit the content of each of these elements to 0.03% or less.
  • the steel structure of the high-strength hot-rolled steel sheet of the present invention has martensite and/or lower bainite as the main phase, and the volume fraction of residual ⁇ is less than 3%.
  • Main phase martensite and/or lower bainite
  • the microstructure in order to obtain high strength and excellent toughness and delayed fracture resistance after post-heating, the microstructure is made to have martensite and/or lower bainite as the main phase. If ferrite, pearlite, residual ⁇ , etc. become the main phase, it becomes difficult to achieve both high strength and excellent toughness and delayed fracture resistance after post-heating. Therefore, the steel microstructure is made to have martensite and/or lower bainite as the main phase.
  • the martensite may be either auto-tempered martensite or tempered martensite, but fresh martensite having no carbides inside is excluded.
  • the lower bainite may be tempered lower bainite.
  • the main phase means a phase that occupies 50% or more in terms of area ratio.
  • the area ratio of the main phase is preferably 60% or more, and more preferably 75% or more.
  • martensite may be the main phase
  • lower bainite may be the main phase
  • the total of martensite and lower bainite may be the main phase.
  • the upper limit of the area ratio of the main phase is not particularly limited and may be 100%.
  • the area ratio of the main phase may be, for example, less than 100% or 98% or less.
  • Amount of retained austenite (residual ⁇ ) less than 3% Since retained austenite (residual ⁇ ) is a structure that significantly reduces strength and toughness by transforming into pearlite after post-heating, it is preferable to reduce it as much as possible.
  • the volume fraction of retained ⁇ is allowed to be less than 3%. Therefore, the volume fraction of retained ⁇ is set to less than 3%.
  • the volume fraction of retained ⁇ is preferably less than 2%, and more preferably less than 1%. There is no particular limit on the lower limit of the volume fraction of retained ⁇ , and the volume fraction of retained ⁇ may be 0%.
  • the phases other than martensite, lower bainite, and residual gamma may be one or more of ferrite, pearlite, and upper bainite.
  • the total area ratio of the other phases is preferably 30% or less, and more preferably 25% or less. There is no particular lower limit to the area ratio of the other phases, and the total area ratio of the other phases may be 0%.
  • (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) is 0.300 or more and less than 0.800. It is preferably 0.350 or more. It is also preferably 0.700 or less.
  • the value of (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb) is determined by the method described in the Examples.
  • the growth of the precipitates competes with the precipitation of new TiC, NbC, etc. during post-heating. This appropriately suppresses the precipitation of fine TiC, NbC, etc., and can suppress excessive strength increase and toughness decrease.
  • the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 mass% or more.
  • the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.030 mass% or less. Therefore, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 to 0.030 mass%. Preferably, it is 0.013 mass% or more. Also, it is preferably 0.027 mass% or less.
  • the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is determined by the method described in the Examples.
  • Pole density of ⁇ 110 ⁇ 111> orientation in the surface layer region from the surface to 100 ⁇ m in the center direction of the sheet thickness 1.8 to 5.0
  • the surface layer region from the surface of the steel plate to 100 ⁇ m in the thickness center direction strongly influences the formation of fracture surface during punching or high-speed deformation.
  • the pole density of the ⁇ 110 ⁇ 111> orientation in this region to the range of 1.8 to 5.0, excellent toughness can be obtained after post-heating, and the fracture surface properties of punching become good, and excellent delayed fracture resistance properties can be obtained after post-heating.
  • the pole density of the ⁇ 110 ⁇ 111> orientation needs to be 1.8 or more in the surface layer region from the surface to 100 ⁇ m in the thickness center direction.
  • the pole density of the ⁇ 110 ⁇ 111> orientation is set to 1.8 to 5.0 in the surface layer region from the surface to 100 ⁇ m in the thickness center direction. It is preferably set to 2.0 or more. It is also preferably set to 4.0 or less, and more preferably set to 3.0 or less.
  • the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region extending from the surface to 100 ⁇ m in the sheet thickness center direction is determined by the method described in the examples.
  • the high-strength hot-rolled steel sheet of the present invention is produced by heating a slab having the above-mentioned composition to a temperature range of 1150 to 1300°C, holding the slab in the temperature range for 0.2 to 3.5 hours, and then hot rolling the slab under conditions in which the total reduction in the temperature range of 1080°C or higher is 80 to 90%, the total reduction in the temperature range of 900°C or lower is 20% or more, and the reduction per pass at or below T (°C) calculated by the following formula is 25% or less, followed by allowing the slab to cool for 1.0 s or more, then cooling the slab at a temperature range up to 550°C at an average cooling rate of 50°C/s or more, setting the time from reaching 550°C to the start of quenching to 0.5 to 4.0 s, then quenching the slab to a coiling temperature of 100 to 250°C at a cooling rate of 200°C/s or more, and
  • T(°C) 800+1000[Ti]+2500[Nb]
  • [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
  • the total reduction in the temperature range of 1080° C. or higher is determined from the ratio of the slab thickness before hot rolling to the plate thickness at 1080° C.
  • the total reduction in the temperature range of 900° C. or lower is determined from the ratio of the plate thickness at 900° C. to the final plate thickness.
  • the reduction per pass at or below T (° C.) is determined from the ratio of the plate thickness before and after each pass of rolling at or below T (° C.).
  • the above temperatures are the surface temperatures at the center of the width of the steel plate, and the above average cooling rates and cooling speeds are the average cooling rate and cooling speed at the surface at the center of the width of the steel plate, respectively. Furthermore, unless otherwise specified, the average cooling rate is [(cooling start temperature - cooling stop temperature) / cooling time from cooling start temperature to cooling stop temperature].
  • the heating temperature of the slab is set to 1150 to 1300°C.
  • the heating temperature is preferably 1170° C. or higher, and more preferably 1185° C. or higher.
  • the heating temperature is preferably 1280° C. or lower, and more preferably 1265° C. or lower.
  • Holding time in the temperature range of 1150 to 1300°C 0.2 to 3.5 hours If the holding time in the temperature range of 1150 to 1300°C is less than 0.2 hours, the dissolution of Ti-containing precipitates and Nb-containing precipitates will be insufficient. As a result, a value of (amount of dissolved Ti + amount of dissolved Nb)/(total amount of Ti + total amount of Nb) of 0.300 or more and less than 0.800, or a value of 0.010 to 0.030 mass% of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more will not be obtained.
  • the holding time of the slab in the above temperature range is set to 0.2 to 3.5 hours.
  • the holding time is preferably 0.4 hours or more.
  • the holding time is preferably 2.5 hours or less.
  • Total reduction rate at temperatures above 1080°C: 80-90% By carrying out a total reduction of 80 to 90% in a temperature range of 1080°C or more, it is possible to promote the generation and growth of coarse Ti-containing precipitates and Nb-containing precipitates having a particle size of 100 nm or more. As a result, the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more can be set to 0.010 to 0.030 mass%. If the total reduction is less than 80%, the generation of precipitates having a particle size of 100 nm or more is insufficient, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more is less than 0.010 mass%.
  • the total reduction in a temperature range of 1080°C or more is set to 80 to 90%.
  • the total reduction is preferably set to 81% or more.
  • the total rolling reduction is preferably 88% or less.
  • Total rolling reduction in the temperature range of 900°C or less is 20% or more If the total rolling reduction in the temperature range of 900°C or less is less than 20%, strain-induced precipitation is suppressed, Ti-containing precipitates and Nb-containing precipitates are reduced, and a value of (solubilized Ti amount + solid-solubilized Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. Or, the texture of the surface layer portion is insufficiently developed, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region is not able to be 1.8 to 5.0. Therefore, the total rolling reduction in the temperature range of 900°C or less is set to 20% or more.
  • the upper limit of the total rolling reduction is not particularly limited, but the total rolling reduction is preferably 80% or less, more preferably 60% or less.
  • T (°C) or less When reduction of more than 25% per pass is applied at T (°C) or less calculated by the following formula, strain-induced precipitation is promoted, Ti-containing precipitates and Nb-containing precipitates increase, and a value of (solute Ti amount + solute Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained.
  • the texture of the surface layer part develops, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. Therefore, the reduction rate per pass at T (°C) or less is set to 25% or less.
  • the reduction rate is preferably set to 20% or less, more preferably set to 18% or less.
  • the lower limit of the reduction rate is not particularly limited, but since coarse grains may occur at 5% or less, the reduction rate is preferably set to more than 5%.
  • the reduction rate is more preferably set to 7% or more.
  • [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
  • Cooling for 1.0 s or more By cooling after rolling under the above conditions, partial strain is released, strain-induced precipitation and dislocation precipitation during subsequent cooling are suppressed, and Ti-containing precipitates and Nb-containing precipitates can be reduced. To obtain such an effect, it is necessary to set the cooling time after rolling to 1.0 s or more.
  • the cooling time is preferably 1.5 s or more, more preferably 2.0 s or more, and even more preferably 2.2 s or more.
  • There is no particular limit to the upper limit of the cooling time but if the cooling time is 5.0 s or less, it becomes easier to control the subsequent hot rolling, so the cooling time is preferably 5.0 s or less.
  • cooling means exposure to the atmosphere (air cooling) without active cooling (accelerated cooling) by water injection or the like.
  • hot rolling includes rough rolling and finish rolling, and the cooling time after rolling is the cooling time after hot rolling, i.e., after finish rolling.
  • Cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C After the above cooling, cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C. If the average cooling rate up to 550°C is less than 50°C/s, excessive generation of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc., and formation of the crystal orientation in the surface layer region will be caused. As a result, the phase structure of the present invention, the precipitates, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 will not be obtained.
  • the average cooling rate in the temperature range from the cooling start temperature to 550°C after the above cooling is set to 50°C/s or more.
  • the average cooling rate is preferably set to 70°C/s or more.
  • the average cooling rate is preferably less than 500°C/s, and more preferably less than 200°C/s.
  • Time from reaching 550°C to starting quenching 0.5 to 4.0 s
  • a certain time leaving a certain time interval
  • starting quenching quenching at a cooling rate of 200°C/s or more, which will be described later
  • bainite can be formed in the medium temperature region near the surface layer.
  • the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of the present invention can be obtained. If the time from reaching 550°C to starting quenching is less than 0.5 s, such an effect cannot be obtained sufficiently, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained.
  • the time from 550°C to starting quenching is set to 0.5 to 4.0 s.
  • the time is preferably 0.7 s or more.
  • the time is preferably 2.0 s or less, more preferably 1.6 s or less.
  • Cooling rate to coiling temperature of 100 to 250°C: 200°C/s or more quenching is started after a time of 0.5 to 4.0 s has elapsed between reaching 550°C and starting quenching. If the cooling (quenching) rate to coiling temperature of 100 to 250°C is less than 200°C/s, upper bainite and residual ⁇ are excessively generated, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region is increased. As a result, the phase structure of the present invention and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region cannot be obtained. Therefore, the cooling rate to the coiling temperature is set to 200°C/s or more. The cooling rate is preferably set to 250°C/s or more. The upper limit of the cooling rate is not particularly limited, but from the viewpoint of shape stability, the cooling rate is preferably 1000°C/s or less, and more preferably 500°C/s or less.
  • Winding temperature 100 to 250°C
  • the coiling temperature is set to 100 to 250°C.
  • the coiling temperature is preferably 120°C or higher.
  • the coiling temperature is preferably 220°C or lower.
  • the high-strength hot-rolled steel sheet of the present invention has excellent strength, toughness, and delayed fracture resistance after post-heating.
  • the heating temperature of the post-heating is 400°C or higher.
  • the upper limit of the heating temperature of the post-heating is not particularly limited, but an example of the heating temperature of the post-heating is 1150°C or lower.
  • the heating time of the post-heating (holding time at the heating temperature) is not particularly limited, but an example of the heating time is more than 0 seconds.
  • the heating time of the post-heating is 3600 seconds or less, for example.
  • the hot-rolled steel sheet was used to observe the structure, analyze solute Ti, solute Nb, Ti-containing precipitates, and Nb-containing precipitates, and evaluate the tensile properties according to the following test methods. Furthermore, the hot-rolled steel sheet was post-heated as shown in Table 2, and the hardness, toughness, and delayed fracture resistance were evaluated according to the following test methods using the hot-rolled steel sheet after post-heating.
  • the post-heating temperature was 400°C or higher, at which an improvement in stretch flangeability is observed, and the post-heating time was 3600 s or less from the viewpoint of productivity.
  • the area ratio of martensite and lower bainite refers to the ratio of the area of each structure to the observed area.
  • the area ratio of martensite was measured by cutting a sample from the obtained hot-rolled steel sheet, polishing the plate thickness cross section parallel to the rolling direction, corroding it with 3% nital, and photographing the plate thickness 1/4 position at a magnification of 1500 times with a SEM (scanning electron microscope) in three fields of view.
  • the area ratio of each structure was calculated from the image data of the obtained secondary electron image using Image-Pro manufactured by Media Cybernetics, and the average area ratio of the three fields of view was taken as the area ratio of each structure.
  • the structure may be determined by a general classification, but can be determined, for example, as follows.
  • lower bainite is distinguished as black or dark gray, gray, or light gray containing oriented carbides.
  • Martensite is a structure of black to light gray that is regular but contains carbides of multiple orientations. Alternatively, it is observed as white or light gray without carbides.
  • the retained austenite is observed as white or light gray without containing carbides. Since it may be difficult to distinguish between a part of martensite and the retained austenite, the retained austenite was obtained by the method described below, and the area ratio of the martensite was obtained by subtracting it from the total area ratio of the martensite and the retained austenite obtained from the SEM image.
  • the carbides are white dots or lines.
  • ferrite is a structure that is black or dark gray and does not have a substructure such as carbides or laths inside
  • pearlite can be distinguished as a black and white layered or partially interrupted layered structure.
  • upper bainite can be distinguished as a structure that is black or dark gray and has a substructure such as carbides or laths inside.
  • the amount of retained ⁇ is obtained as follows.
  • the hot-rolled steel sheet was ground to 1/4+0.1 mm of the sheet thickness, and then chemically polished to a further 0.1 mm to obtain the measurement surface.
  • the measurement surface was measured using an X-ray diffraction apparatus with Mo K ⁇ 1 radiation to measure the integral reflection intensities of the (200), (220), and (311) surfaces of fcc iron (austenite) and the (200), (211), and (220) surfaces of bcc iron (ferrite).
  • the volume fraction was then calculated from the intensity ratio of the integral reflection intensity from each surface of the fcc iron to the integral reflection intensity from each surface of the bcc iron, and this was taken as the amount of residual ⁇ .
  • Table 3 shows the structures constituting the main phase and other structures that account for 50% or more of the area ratio of each structure obtained.
  • M means martensite
  • LB means lower bainite
  • ⁇ means retained austenite
  • O means other phases.
  • the other phases include one or more of ferrite, pearlite, and upper bainite.
  • Pole density of ⁇ 110 ⁇ 111> orientation in the surface region from the surface to 100 ⁇ m in the thickness center direction Samples were cut out from the obtained hot-rolled steel sheet, the thickness cross section parallel to the rolling direction was polished, and strain was removed by electrolytic polishing. After that, crystal orientation data was obtained for the surface region from the surface to 100 ⁇ m in the thickness center direction by EBSD (electron backscatter diffraction). Each sample was measured in three fields of view with a measurement area of 100 ⁇ m x 100 ⁇ m, an acceleration voltage of 30 kV, and a step size of 100 nm. OIM Analysis Ver. 7.3.0 manufactured by TSL Solutions was used to analyze the obtained data.
  • the average value of the pole densities of the three fields of view of each sample was taken as the pole density of each sample.
  • the concentration ratios of Ti and Nb to Fe were calculated, and further multiplied by the content (mass%) of Fe in the test piece to obtain the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%).
  • the content (mass%) of Fe in the test piece was obtained by subtracting the total content (mass%) of components other than Fe from 100 mass%.
  • the ratio of the total dissolved Ti amount (mass%) and dissolved Nb amount (mass%) to the total contained Ti amount (mass%) and contained Nb amount (mass%) was calculated.
  • the test piece having the precipitate attached to the surface after electrolysis was taken out from the electrolytic solution and immersed in an aqueous solution of sodium hexametaphosphate (500 mg/L) (hereinafter referred to as an aqueous SHMP solution). Then, ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution. Next, the aqueous SHMP solution containing the precipitate was filtered using a filter with a pore size of 100 nm, and the precipitate collected on the 100 nm filter was decomposed with acid, and the decomposition solution was analyzed using an ICP emission spectrometer, and the absolute values of Ti and Nb in the decomposition solution were measured.
  • an aqueous SHMP solution sodium hexametaphosphate (500 mg/L)
  • ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution.
  • the absolute values of Ti and Nb obtained were divided by the amount of electrolyte to obtain the amount of Ti and the amount of Nb (mass %) contained in the precipitates having a particle size of 100 nm or more when the total composition of the test piece was taken as 100 mass %).
  • the total of the obtained amount of Ti (mass %) and amount of Nb (mass %) was divided by the total amount of Ti (mass %) and amount of Nb (mass %) contained in the test piece to obtain the total amount of Ti (mass %) present as precipitates containing Ti having a particle size of 100 nm or more and the amount of Nb (mass %) present as precipitates containing Nb having a particle size of 100 nm or more.
  • the amount of electrolyte was obtained by measuring the mass of the test piece after the precipitates were peeled off and subtracting it from the mass of the test piece before electrolysis.
  • JIS No. 5 tensile test pieces (JIS Z 2241:2011) were taken from the obtained hot-rolled steel sheets in the direction parallel to the rolling direction, and a tensile test was carried out in accordance with the provisions of JIS Z 2241:2011 at a strain rate of 10 ⁇ 3 /s to determine TS. In the present invention, a TS of 1180 MPa or more was considered to be acceptable.
  • Vickers hardness test Samples were cut out from the obtained hot-rolled steel sheet and the hot-rolled steel sheet after post-heating, and the cross section of the sheet thickness parallel to the rolling direction was polished. Then, a Vickers hardness test was performed at 1/4 of the sheet thickness position with a load of 5 kg and five measurement points, and the average (arithmetic mean) was taken as the Vickers hardness of the steel sheet. A difference in hardness ( ⁇ HV) of 50 or less before and after post-heating was judged to be excellent in strength after post-heating and was considered to have passed the test.
  • ⁇ HV hardness
  • Charpy impact test From the hot-rolled steel sheet obtained by post-heating the hot-rolled steel sheet, a test piece with a width of 10 mm and a length of 55 mm was taken, and a V-notch with a tip angle of 45°, a tip radius of 0.25 mm, and a depth of 2 mm was made to prepare a Charpy impact test piece. Then, in accordance with JIS Z 2242:2018, a Charpy impact test was performed five times at -20 ° C. to evaluate the ductile fracture rate. A test piece with an average ductile fracture rate of 50% or more after five tests was judged to have excellent toughness after post-heating and was passed. The plate thickness was 2.9 mm, and the notch direction was parallel to the rolling direction.
  • Delayed fracture test From the obtained hot-rolled steel sheet, a test piece with a width of 30 mm and a length of 110 mm was taken, and the post-heating treatment shown in Table 2 was performed to obtain a test piece. This was subjected to 90° V-bending with a bending radius of 15 mm so that the ridge line was parallel to the rolling direction, and bolts were tightened by the amount of opening due to springback, and the test piece was immersed in hydrochloric acid of pH 3 for 96 hours to check for the presence or absence of cracks. Those that did not have cracks were judged to have excellent delayed fracture resistance after post-heating and were passed.
  • the end faces of the test pieces were formed by shearing with a shear angle of 1° and a clearance of 10%, and burrs were bent on the outside.
  • the "delayed fracture time (hr)” in Table 3 indicates the time when cracks occurred in the test piece. However, “96” in the “delayed fracture time (hr)” indicates that no cracks occurred in the test piece after the above 96-hr immersion.
  • All of the inventive examples have a TS of 1180 MPa or more, and are excellent in strength, toughness, and delayed fracture resistance after post-heating.
  • the comparative examples that fall outside the scope of the present invention either do not have the desired strength (TS) or do not achieve one or more of the desired strength, toughness, and delayed fracture resistance after post-heating.
  • the present invention it is possible to obtain a high-strength hot-rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and excellent strength, toughness and delayed fracture resistance after post-heating.
  • the high-strength steel sheet of the present invention is used for automobile parts, it can greatly contribute to improving the collision safety and fuel efficiency of automobiles.

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Abstract

L'invention fournit une tôle en acier laminée à chaud hautement résistante excellente en termes de résistance après postchauffage, de ténacité et de résistance à la rupture différée.　Plus précisément, l'invention concerne une tôle en acier laminée à chaud hautement résistante qui possède une composition contenant, en % en masse, ０,０６ à ０,２３％ de C, ０,１ à ３,０％ de Si, １,５ à ３,５％ de Mn, plus de ０％ à ０,０５０％ ou moins de P, plus de ０％ à ０,００５０％ ou moins de S, plus de ０％ à １,５％ ou moins de Al, plus de ０％ à ０,０１０％ ou moins de N et plus de ０％ à ０,００３％ ou moins de O, le reste étant constitué de Fe ou des impuretés inévitables. En outre, la tôle en acier de l'invention présente une structure d'acier telle qu'une martensite et/ou une bainite inférieure constitue sa phase principale, une austénite résiduelle représente moins de 3% en volume, （quantité de Ti en solution solide ＋ quantité de Nb en solution solide）／（quantité de Ti totale ＋ quantité de Nb totale） est supérieur ou égal à ０,３００ et inférieur à ０,８００, au total la quantité de Ti et la quantité de Nb présents en tant que dépôt de diamètre particulaire supérieur ou égal à 100nm, sont comprises entre ０,０１０ et ０,０３０% en masse, et la densité polaire d'une orientation de ｛１１０｝＜１１１＞ dans une région de couche superficielle, est comprise entre １,８ et ５,０.
PCT/JP2023/041535 2022-11-22 2023-11-17 Tôle en acier laminée à chaud hautement résistante, et procédé de fabrication de celle-ci WO2024111525A1 (fr)

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Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2014132968A1 (fr) * 2013-02-26 2014-09-04 新日鐵住金株式会社 TÔLE D'ACIER LAMINÉE À CHAUD À HAUTE RÉSISTANCE, DOTÉE D'UNE RÉSISTANCE À LA TRACTION MAXIMALE DE 980 MPa OU SUPÉRIEURE ET PRÉSENTANT D'EXCELLENTES TREMPABILITÉ PAR CUISSON ET TÉNACITÉ À BASSES TEMPÉRATURES
WO2014185405A1 (fr) * 2013-05-14 2014-11-20 新日鐵住金株式会社 Feuille d'acier laminee a chaud et son procede de production
WO2014188966A1 (fr) * 2013-05-21 2014-11-27 新日鐵住金株式会社 Tôle d'acier laminée à chaud et son procédé de fabrication
WO2018138887A1 (fr) * 2017-01-27 2018-08-02 新日鐵住金株式会社 Tôle d'acier et tôle d'acier plaquée
WO2022042729A1 (fr) * 2020-08-31 2022-03-03 宝山钢铁股份有限公司 Acier à expansion de trou ultra-élevée, à martensite à très faible teneur en carbone et à austénite résiduelle de qualité 980 mpa, et son procédé de fabrication
WO2022042730A1 (fr) * 2020-08-31 2022-03-03 宝山钢铁股份有限公司 Acier à forte expansion de trou martensitique à faible teneur en carbone et à haute résistance et procédé de fabrication

Patent Citations (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
WO2014132968A1 (fr) * 2013-02-26 2014-09-04 新日鐵住金株式会社 TÔLE D'ACIER LAMINÉE À CHAUD À HAUTE RÉSISTANCE, DOTÉE D'UNE RÉSISTANCE À LA TRACTION MAXIMALE DE 980 MPa OU SUPÉRIEURE ET PRÉSENTANT D'EXCELLENTES TREMPABILITÉ PAR CUISSON ET TÉNACITÉ À BASSES TEMPÉRATURES
WO2014185405A1 (fr) * 2013-05-14 2014-11-20 新日鐵住金株式会社 Feuille d'acier laminee a chaud et son procede de production
WO2014188966A1 (fr) * 2013-05-21 2014-11-27 新日鐵住金株式会社 Tôle d'acier laminée à chaud et son procédé de fabrication
WO2018138887A1 (fr) * 2017-01-27 2018-08-02 新日鐵住金株式会社 Tôle d'acier et tôle d'acier plaquée
WO2022042729A1 (fr) * 2020-08-31 2022-03-03 宝山钢铁股份有限公司 Acier à expansion de trou ultra-élevée, à martensite à très faible teneur en carbone et à austénite résiduelle de qualité 980 mpa, et son procédé de fabrication
WO2022042730A1 (fr) * 2020-08-31 2022-03-03 宝山钢铁股份有限公司 Acier à forte expansion de trou martensitique à faible teneur en carbone et à haute résistance et procédé de fabrication

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