WO2016147549A1 - Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication - Google Patents

Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication Download PDF

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WO2016147549A1
WO2016147549A1 PCT/JP2016/000778 JP2016000778W WO2016147549A1 WO 2016147549 A1 WO2016147549 A1 WO 2016147549A1 JP 2016000778 W JP2016000778 W JP 2016000778W WO 2016147549 A1 WO2016147549 A1 WO 2016147549A1
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steel sheet
temperature
rolled steel
strength
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PCT/JP2016/000778
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Japanese (ja)
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真平 吉岡
克利 ▲高▼島
長谷川　浩平
船川　義正
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Ｊｆｅスチール株式会社
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Priority to JP2017506047A priority Critical patent/JP6210175B2/ja
Priority to CN201680016385.8A priority patent/CN107429344A/zh
Priority to MX2017011825A priority patent/MX2017011825A/es
Priority to US15/556,448 priority patent/US20180037969A1/en
Priority to EP16764383.2A priority patent/EP3272892B1/fr
Priority to KR1020177025309A priority patent/KR101998652B1/ko
Publication of WO2016147549A1 publication Critical patent/WO2016147549A1/fr

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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
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    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/001—Heat treatment of ferrous alloys containing Ni
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D6/00—Heat treatment of ferrous alloys
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    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
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    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
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    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/16—Ferrous alloys, e.g. steel alloys containing copper
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    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • 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
    • C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of 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
    • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
    • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
    • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
    • C23G1/02—Cleaning or pickling metallic material with solutions or molten salts with acid solutions
    • C23G1/08—Iron or steel
    • C—CHEMISTRY; METALLURGY
    • C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
    • C23G—CLEANING OR DE-GREASING OF METALLIC MATERIAL BY CHEMICAL METHODS OTHER THAN ELECTROLYSIS
    • C23G1/00—Cleaning or pickling metallic material with solutions or molten salts
    • C23G1/02—Cleaning or pickling metallic material with solutions or molten salts with acid solutions
    • C23G1/08—Iron or steel
    • C23G1/081—Iron or steel solutions containing H2SO4
    • 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
    • C21D2211/001—Austenite
    • 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
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing

Definitions

  • the present invention relates to a high-strength cold-rolled steel sheet that has a tensile strength (TS) of 1300 MPa or more and is excellent in chemical conversion treatment and workability, and a method for producing the same, useful for the use of automobile members.
  • TS tensile strength
  • automotive steel sheets are used after being coated, and chemical conversion treatment such as phosphate treatment is performed as a pretreatment for the coating. Since the chemical conversion treatment of this steel plate is one of the important treatments for ensuring the corrosion resistance after coating, the automotive steel plate is also required to have excellent chemical conversion treatment properties.
  • Patent Document 1 the balance between strength and ductility is improved by adding a large amount of C.
  • a large amount of C is added, stretch flangeability is deteriorated due to the hardness difference between the two phases.
  • Patent Document 2 Si is utilized. However, when a large amount of Si is added, in the case of the production method described in Patent Document 2, it is speculated that Si oxide is formed on the surface of the steel sheet in the continuous annealing line, and the chemical conversion property is deteriorated. It is not preferable.
  • Patent Document 3 by adding a large amount of Mn, the Si-Mn composite oxide is finely dispersed on the steel sheet surface and used as a nucleation site for zinc phosphate crystals, and the steel sheet surface SiO 2 is reduced as much as possible. To ensure chemical conversion. However, it is difficult to achieve a tensile strength of 1300 MPa and an elongation of 10% or more with the amounts of C and Si described in Patent Document 3.
  • an object of the present invention is to provide a high-strength cold-rolled steel sheet having a tensile strength of 1300 MPa or more and excellent in chemical conversion property and workability, and a method for producing the same.
  • the microstructure needs to be a martensite single phase structure or a ferrite-martensite composite structure.
  • optimization of component design, structure control, etc. is important to achieve both high strength and workability.
  • Mn is effective for increasing the strength of steel sheets.
  • Mn is added more than necessary, it segregates during casting, and a steel structure in which ferrite and martensite are distributed in a band shape is formed. For this reason, anisotropy occurs in the mechanical characteristics, and the workability deteriorates.
  • Component composition is mass%, C: 0.15% to 0.22%, Si: 1.0% to 2.0%, Mn: 1.7% to 2.5%, P: 0.05% or less, S: 0.02% or less, Al: 0.01% or more and 0.05% or less, N: 0.005% or less, satisfying the following formula (1), the balance is composed of iron and inevitable impurities, the structure is area ratio, tempered martensite is 60% More than 100%, untransformed austenite is 5% or less (including 0%), the balance is ferrite, and the average crystal grain size of the ferrite is less than 3.5 ⁇ m.
  • [Si] represents the Si content (mass%)
  • [Mn] represents the Mn content (mass%).
  • the composition contains one or more of% by mass: V: 0.01% to 0.30%, Mo: 0.01% to 0.30%, Cr: 0.01% to 0.30% [1 ]
  • Cu includes 0.01% or more and 0.30% or less, and Ni: 0.01% or more and 0.30% or less in mass%.
  • Sn 0.001% to 0.100%
  • Sb 0.001% to 0.100%
  • Ca 0.0002% to 0.0100%
  • W 0.01% to 0.10%
  • the high-strength cold-rolled steel sheet according to any one of the above [1] to [6] which contains at least one of Co: 0.01% to 0.10% and REM: 0.0002% to 0.0050%.
  • a steel material having the composition described in any one of [1] to [7] above is heated to a temperature of 1200 ° C. or higher, and then hot rolled to a finish rolling outlet temperature of 800 ° C. or higher. Winding at a temperature of 450 ° C to 700 ° C, cold rolling, then heating to an annealing temperature of Ac 1 point or more and Ac 3 point or less, and the residence time in the temperature range from Ac 1 point to Ac 3 point 30 seconds or more and 1200 seconds or less, the primary cooling from the annealing temperature to the primary cooling stop temperature of 600 ° C or higher at an average cooling rate of less than 100 ° C / s, and the secondary cooling stop temperature of 100 ° C or lower to the average cooling rate An annealing treatment is performed to perform secondary cooling at 100 ° C./s or more and 1000 ° C./s or less, and then heating is performed to a temperature of 100 ° C.
  • the high-strength cold-rolled steel sheet is a cold-rolled steel sheet having a tensile strength (TS) of 1300 MPa or more.
  • a high-strength cold-rolled steel sheet having a tensile strength of 1300 MPa or more and excellent in chemical conversion property and workability can be obtained.
  • the high-strength cold-rolled steel sheet of the present invention has a tensile strength of 1300 MPa or more and is excellent in chemical conversion treatment and workability, so it can be suitably used for the use of automobile structural members, etc. The effect is remarkable, such as improving its reliability.
  • C 0.15% to 0.22%
  • C is an element effective for improving the balance between strength and ductility of the steel sheet. If the C content is less than 0.15%, it is difficult to ensure a tensile strength of 1300 MPa or more. On the other hand, when the C content exceeds 0.22%, coarse cementite precipitates and workability such as stretch flangeability deteriorates. Therefore, the C content is in the range of 0.15% to 0.22%. Preferably it is 0.16% or more. Preferably it is 0.20% or less.
  • Si 1.0% or more and 2.0% or less Si is an element effective for ensuring strength without significantly reducing the ductility of the steel sheet.
  • the Si content is less than 1.0%, a steel plate with high strength and high workability cannot be produced.
  • the amount of Si exceeds 2.0%, even if a step of re-acid washing after pickling is performed, the Si oxide on the surface of the steel sheet cannot be removed, and the chemical conversion treatment performance is lowered. Therefore, the Si content is in the range of 1.0% to 2.0%. Preferably it is 1.0% or more. Preferably it is 1.5% or less.
  • Mn 1.7% to 2.5%
  • Mn is an element that increases the strength of the steel sheet. If the Mn content is less than 1.7%, it is difficult to ensure a tensile strength of 1300 MPa or more. On the other hand, when the amount of Mn exceeds 2.5%, a steel structure in which ferrite and martensite are distributed in a band shape is formed due to segregation during casting. As a result, anisotropy occurs in the mechanical characteristics, and workability deteriorates. Therefore, the Mn content is in the range of 1.7% to 2.5%.
  • [Si] represents the Si content (mass%)
  • [Mn] represents the Mn content (mass%)
  • the amount of Si-based oxide and Si-Mn composite oxide is determined by the balance between Si and Mn.
  • Si silicon-based oxide
  • Si-Mn composite oxide an oxide mainly composed of Si-Mn (Si-Mn composite oxide) is generated excessively, and the present invention.
  • the intended chemical conversion processability cannot be obtained. Therefore, [Si] / [Mn] ⁇ 0.5.
  • P 0.05% or less
  • P is an impurity element and must be reduced to degrade ductility. If it exceeds 0.05%, the local ductility deteriorates due to grain boundary embrittlement accompanying P segregation to the austenite grain boundaries during casting. As a result, the balance between strength and ductility deteriorates. Therefore, the P content is 0.05% or less. Preferably it is 0.02% or less.
  • S 0.02% or less S is present as MnS in the steel sheet, and causes a reduction in impact resistance, strength, and stretch flangeability. Therefore, the upper limit is 0.02%. Preferably it is 0.002% or less.
  • Al 0.01% or more and 0.05% or less Al has an effect of reducing ductility by reducing oxides such as Si by itself forming an oxide. However, a significant effect cannot be obtained at less than 0.01%.
  • Al is excessively added exceeding 0.05%, Al and N are combined to form a nitride. Since this nitride precipitates on the austenite grain boundary during casting and embrittles the grain boundary, it deteriorates stretch flangeability. Therefore, the Al content is in the range of 0.01% to 0.05%.
  • N forms nitrides with Al and Ti and deteriorates stretch flangeability as described above.
  • the N content exceeds 0.005%, the stretch flangeability is remarkably deteriorated by Ti and Al nitrides, and the decrease in elongation due to the increase in solute N is also remarkable. Therefore, the N content is 0.005% or less. Preferably it is 0.002% or less.
  • Ti 0.010% or more and 0.020% or less Ti has an effect of refining the structure, and may be added as necessary. If the amount of Ti is less than 0.010%, the effect of refining the structure is small. On the other hand, even if added over 0.020%, not only the effect of refining the structure is saturated, but also coarse Ti and Nb composite carbides may be formed to deteriorate the balance between strength and ductility and stretch flangeability. Further, the manufacturing cost increases. For this reason, when adding Ti, it is made into 0.010% or more and 0.020% or less. Preferably it is 0.012% or more. Preferably it is 0.018% or less.
  • Nb 0.02% or more and 0.10% or less Nb has the effect of refining the structure in the same manner as Ti, and may be added as necessary. If the Nb content is less than 0.02%, the effect of refining the structure is small. On the other hand, adding over 0.10% not only saturates the effect of refining the structure but also forms coarse Ti and Nb composite carbides, which may deteriorate the balance between strength and ductility and stretch flangeability. Furthermore, the manufacturing cost increases. For this reason, when adding Nb, it is made into 0.02% or more and 0.10% or less. Preferably it is 0.04% or more. Preferably it is 0.08% or less.
  • B 0.0002% or more and 0.0020% or less B segregates at the austenite grain boundaries during heating in continuous annealing, suppresses ferrite transformation and bainite transformation from austenite during cooling, and facilitates the formation of tempered martensite. As a result, the steel plate is strengthened. Therefore, you may add as needed. If the amount of B is less than 0.0002%, the above effect is small. On the other hand, if the amount of B exceeds 0.0020%, borocarbide Fe 23 (C, B) 6 is generated, which may cause deterioration of workability and strength. For this reason, the B content is 0.0002% or more and 0.0020% or less when added.
  • V 0.01% to 0.30%
  • Mo 0.01% to 0.30%
  • Cr 0.01% to 0.30%
  • V 0.01% or more and 0.30% or less Fine carbide formed by combining V and C is effective for precipitation strengthening of the steel sheet, and V may be added as necessary. The effect is small when the V content is less than 0.01%. On the other hand, if the amount of V exceeds 0.30%, carbides may precipitate excessively and the balance between strength and ductility may deteriorate. For this reason, the amount of V is 0.01% or more and 0.30% or less when added.
  • Mo 0.01% or more and 0.30% or less Mo is effective for strengthening the quenching of the steel sheet, and has the effect of refining the steel structure, so it may be added as necessary. The effect is small when the Mo content is less than 0.01%. On the other hand, if the amount of Mo exceeds 0.30%, not only the effect is saturated, but also the formation of Mo oxide on the steel sheet surface is promoted during continuous annealing, and the chemical conversion treatment property of the steel sheet may be remarkably lowered. For this reason, the Mo amount is 0.01% or more and 0.30% or less when added.
  • Cr 0.01% or more and 0.30% or less Cr is effective for strengthening the quenching of the steel sheet, and may be added as necessary. If the Cr content is less than 0.01%, the strengthening ability is small. On the other hand, if the Cr content exceeds 0.30%, the formation of Cr oxides on the steel sheet surface is promoted during continuous annealing, so the chemical conversion property of the steel sheet may be significantly reduced. For this reason, when adding Cr, it is 0.01% or more and 0.30% or less. In the present invention, when further improving the characteristics, it is preferable to contain one or more of Cu: 0.01% to 0.30% and Ni: 0.01% to 0.30%.
  • Cu 0.01% or more and 0.30% or less
  • Cu suppresses ferrite transformation and bainite transformation from austenite during cooling in continuous annealing, facilitates the formation of tempered martensite, and strengthens the steel sheet. Therefore, you may add as needed. If the amount of Cu is less than 0.01%, the above effect is small. On the other hand, if the amount of Cu exceeds 0.30%, ferrite transformation is excessively suppressed and ductility may be reduced. For this reason, when adding Cu, it is 0.01% or more and 0.30% or less.
  • Ni 0.01% or more and 0.30% or less Ni suppresses ferrite transformation and bainite transformation from austenite during cooling in continuous annealing, facilitates the formation of tempered martensite, and strengthens the steel sheet. Therefore, you may add as needed.
  • the amount of Ni is less than 0.01%, the above effect is small.
  • the Ni content exceeds 0.30%, ferrite transformation is excessively suppressed and ductility may be reduced. For this reason, when Ni is added, it is set to 0.01% or more and 0.30% or less.
  • Sn 0.001% to 0.100%
  • Sb 0.001% to 0.100%
  • Ca 0.0002% to 0.0100%
  • W It is preferable to contain one or more of 0.01% to 0.10%, Co: 0.01% to 0.10%, REM: 0.0002% to 0.0050%.
  • Sn 0.001% or more and 0.100% or less
  • Sb 0.001% or more and 0.100% or less Since Sn and Sb have the effect of suppressing surface oxidation, decarburization, and nitriding, they can be contained as necessary. However, the above effects are small when the Sn content and the Sb content are each less than 0.001%. On the other hand, even if the added amount exceeds 0.100%, the effect is saturated. For this reason, when adding Sn and Sb, it is made 0.001% or more and 0.100% or less, respectively. Preferably it is 0.005% or more. Preferably it is 0.010% or less.
  • Ca 0.0002% or more and 0.0100% or less
  • Ca has an effect of improving ductility through the form control of sulfide, grain boundary strengthening, and solid solution strengthening, and can be contained as necessary.
  • the effect is small when the Ca content is less than 0.0002%.
  • ductility will deteriorate by grain boundary segregation. For this reason, when adding Ca, it is made into 0.0002% or more and 0.0100% or less.
  • W 0.01% or more and 0.10% or less
  • Co 0.01% or more and 0.10% or less
  • W and Co have the effect of improving ductility through sulfide morphology control, grain boundary strengthening, and solid solution strengthening. It can be included. However, the effect is small when the W content and the Co content are each less than 0.01%. On the other hand, if added excessively, ductility deteriorates due to grain boundary segregation and the like. For this reason, when adding W and Co, they are 0.01% or more and 0.10% or less, respectively.
  • REM 0.0002% or more and 0.0050% or less REM has an effect of improving ductility through sulfide morphology control, grain boundary strengthening, and solid solution strengthening, and can be contained as necessary.
  • the REM amount is less than 0.0002%, the above effect is small.
  • ductility deteriorates due to grain boundary segregation and the like. For this reason, when adding REM, it is made into 0.0002% or more and 0.0050% or less.
  • the balance other than the above is Fe and inevitable impurities.
  • Inevitable impurities include O (oxygen) and the like, and an O content of 0.01% or less is acceptable.
  • tempered martensite contains 60% or more and less than 100%, untransformed austenite is 5% or less (including 0%), the balance is ferrite, and the average grain size of ferrite is less than 3.5 ⁇ m.
  • Tempered martens The tensile strength of steel having a structure including sites and ferrite increases as the area ratio of tempered martensite increases. In tempered martensite and ferrite, tempered martensite has higher hardness, and deformation resistance during tensile deformation is borne by tempered martensite, which is a hard phase. The larger the tempered martensite area ratio, the more tempered martensite. This is because it approaches the tensile strength of the single phase structure.
  • the area ratio of tempered martensite is less than 40%, a tensile strength of 1300 MPa or more cannot be obtained.
  • the area of the interface between tempered martensite and ferrite is large, that is, when the area ratio of tempered martensite is 40% or more and less than 60%, the frequency of void formation due to the hardness difference between the two phases increases and the voids are connected. It becomes easy to do, and the progress of a crack is accelerated, and stretch flangeability will deteriorate. From the above, in order to improve workability while ensuring tensile strength, the area ratio of tempered martensite needs to be 60% or more.
  • the area ratio of tempered martensite is 100%, excellent workability cannot be obtained. In addition, 5% or less of untransformed austenite may be inevitably mixed. However, if it is 5% or less, there is no problem in obtaining the effect of the present invention and it is allowed. From the above, the area ratio of tempered martensite is less than 100%, untransformed austenite is 5% or less (including 0%), and the balance is ferrite.
  • the lower limit of the area ratio of a suitable tempered martensite is 70%.
  • a preferred upper limit is 90%.
  • the average crystal grain size of ferrite is 3.5 ⁇ m or more, the predetermined strength cannot be obtained because the grain refinement strengthening is insufficient. In addition, since deformation is likely to occur between crystal grains during deformation, workability deteriorates. Therefore, the average crystal grain size of ferrite is less than 3.5 ⁇ m.
  • the area ratio of tempered martensite, the area ratio of ferrite, and the average crystal grain size of ferrite can be measured by the methods of Examples described later.
  • the number of Si-Mn composite oxides having an equivalent circle diameter of 5 ⁇ m or less is less than 10/100 ⁇ m 2 , the presence of Si-Mn composite oxides on the steel sheet surface significantly deteriorates the chemical conversion treatment property. Needless to say, if a coarse Si-Mn composite oxide is present on the surface of the steel sheet, the chemical conversion treatment performance deteriorates. Even in the case of a Si-Mn composite oxide having a circle equivalent diameter of 5 ⁇ m or less, when the distribution form exceeds a certain number density, deterioration of chemical conversion treatment becomes obvious. Therefore, the number of Si-Mn composite oxides having an equivalent circle diameter of 5 ⁇ m or less is defined as less than 10/100 ⁇ m 2 . If it is 10 pieces / 100 ⁇ m 2 or more, a region where zinc phosphate crystals are not formed becomes obvious, and the chemical conversion treatment performance deteriorates. The number is preferably 0/100 ⁇ m 2 .
  • the number of Si-Mn composite oxides having an equivalent circle diameter of 5 ⁇ m or less can be measured by the method of Examples described later.
  • the surface is a range from the surface layer to the position of 3% with respect to the plate thickness in the plate thickness direction.
  • the steel surface coverage of the oxide mainly composed of Si is 1% or less. If an oxide mainly composed of Si is present on the surface of the steel sheet, the chemical conversion treatment performance is remarkably lowered. Therefore, the steel sheet surface coverage of oxide mainly composed of Si is set to 1% or less. Preferably it is 0%.
  • the oxide mainly composed of Si is, for example, SiO 2 . Further, the oxide mainly composed of Si can be measured by the method of Examples described later.
  • the structure, the number of Si-Mn composite oxides, and the steel sheet surface coverage of the oxide mainly composed of Si can be controlled by controlling pickling after annealing, particularly re-acid cleaning, in the manufacturing method described later. Obtainable.
  • the high-strength cold-rolled steel sheet of the present invention heats the steel material (steel slab) having the above-described composition to a temperature of 1200 ° C or higher, and then performs hot rolling to a finish rolling temperature of 800 ° C or higher, and 450 ° C or higher. Winding at a temperature of 700 ° C or less and cold rolling. Next, heating is performed to an annealing temperature of Ac 1 point or more and Ac 3 point or less, and the residence time in the temperature range from Ac 1 point to Ac 3 point is 30 seconds or more and 1200 seconds or less, and the primary cooling stop temperature is 600 ° C. from the annealing temperature.
  • the primary cooling is performed at an average cooling rate of less than 100 ° C./s until the above, and the secondary cooling is performed at an average cooling rate of 100 ° C./s to 1000 ° C./s to a secondary cooling stop temperature of 100 ° C. or lower.
  • it is heated to a temperature of 100 ° C. or higher and 300 ° C. or lower, subjected to a tempering treatment in which the residence time in the temperature range from 100 ° C. to 300 ° C. is 120 seconds or longer and 1800 seconds or shorter, and further pickled and re- pickled.
  • the high-strength cold-rolled steel sheet of the present invention can be manufactured.
  • Ac 1 point and Ac 3 point are values (° C.) obtained from a transformation expansion curve obtained at an average heating rate of 3 ° C./s using a thermal expansion measuring device.
  • the method for melting steel is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed. Further, secondary refining may be performed in a vacuum degassing furnace. Thereafter, the slab (steel material) is preferably formed by a continuous casting method from the viewpoint of productivity and quality, but the slab may be formed by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. .
  • the heating temperature of the steel material is 1200 ° C. or higher and the heating temperature is lower than 1200 ° C., the carbide does not re-dissolve and the workability deteriorates. Therefore, the heating temperature of the steel material is set to 1200 ° C or higher. If the heating temperature becomes too high, it leads to an increase in scale loss accompanying an increase in oxidation mass. Therefore, the heating temperature of the steel material is preferably 1300 ° C. or lower.
  • the heating temperature of the steel material is preferably 1300 ° C. or lower.
  • the rough rolling conditions are not particularly limited.
  • Finishing rolling exit temperature 800 ° C. or more
  • the finish rolling exit temperature is set to 800 ° C. or higher.
  • the upper limit of the finish rolling exit temperature is not particularly limited, but 1000 ° C. or less is preferable because rolling at an excessively high temperature causes scale wrinkles and the like.
  • Winding temperature 450 ° C or more and 700 ° C or less
  • the coiling temperature after hot rolling is lower than 450 ° C, the processed structure generated by hot rolling remains, and the rolling load of the next cold rolling increases.
  • the coiling temperature exceeds 700 ° C., coarse grains are formed, the steel sheet structure becomes non-uniform, and the ductility decreases. Therefore, the coiling temperature is set to 450 ° C. or more and 700 ° C. or less.
  • the lower limit of the suitable winding temperature is 500 ° C.
  • a preferred upper limit is 650 ° C.
  • pickling and then cold rolling as necessary.
  • the conditions for pickling are not particularly limited. Cold rolling needs to be performed to obtain a desired plate thickness. Although there is no restriction
  • annealing temperature residence time 30 seconds or more of a temperature range of from Ac 1 point or more Ac 3 point below the annealing temperature to heating Ac 1 point to Ac 3 point is less than 1 point Ac, predetermined during annealing Austenite (transformed into martensite after quenching) necessary for securing the strength is not generated, and a predetermined strength cannot be obtained even if quenching is performed after annealing.
  • the annealing temperature is more than Ac 3 point, it is possible to obtain martensite with an area ratio of 60% or more by controlling the area ratio of ferrite precipitated during cooling from the annealing temperature. When annealing at more than 3 points, it becomes difficult to obtain a desired metal structure.
  • the annealing temperature is set to Ac 1 point or more and Ac 3 point or less.
  • the annealing temperature is preferably set to 780 ° C. or more from the viewpoint of stably securing an equilibrium area ratio of austenite of 60% or more.
  • the residence time at the annealing temperature is too short, the microstructure is not sufficiently annealed and becomes a non-uniform structure in which a cold-rolled processed structure exists, and the ductility is lowered.
  • the residence time is 30 to 1200 seconds.
  • a preferred lower limit of residence time is 150 seconds.
  • a preferred upper limit is 600 seconds.
  • Secondary cooling stop temperature at an average cooling rate of less than 100 ° C / s from the primary cooling annealing temperature to less than 100 ° C / s with an average cooling rate of less than 100 ° C / s. ) (Slow cooling). It becomes possible to precipitate ferrite during the slow cooling from the annealing temperature, and to control the balance between strength and ductility.
  • the slow cooling stop temperature primary cooling stop temperature
  • the slow cooling stop temperature is less than 600 ° C.
  • a large amount of pearlite is generated in the microstructure and the strength rapidly decreases, so that a tensile strength of 1300 MPa or more cannot be obtained.
  • 680 ° C In order to obtain a predetermined strength more stably, 680 ° C.
  • the average cooling rate is 100 ° C./s or more, a sufficient amount of ferrite does not precipitate during cooling, so that excellent ductility cannot be obtained.
  • the ductility of the metal structure having tempered martensite and ferrite intended in the present invention is attributed to the high work-hardening ability that is manifested by the mixture of hard tempered martensite and soft ferrite.
  • the average cooling rate is 100 ° C./s or more, carbon concentration in the austenite during cooling becomes insufficient, and hard martensite cannot be obtained during rapid cooling. As a result, the work hardening ability of the final structure is lowered and sufficient ductility cannot be obtained.
  • the average cooling rate is less than 100 ° C / s. In order to sufficiently cause carbon concentration in austenite, an average cooling rate of 5 ° C./s or less is preferable.
  • Secondary cooling down temperature to 100 ° C or less Secondary cooling at an average cooling rate of 100 ° C / s or more and 1000 ° C / s or less Following the above-mentioned slow cooling, 100 ° C or less at an average cooling rate of 100 ° C / s or more and 1000 ° C / s or less Cool down (rapidly cool) to the secondary cooling stop temperature. Rapid cooling after slow cooling is performed in order to transform austenite into martensite, but when the average cooling rate is less than 100 ° C / s, austenite transforms into ferrite, bainite or pearlite during cooling, so that a predetermined strength is obtained. I can't.
  • the average cooling rate during rapid cooling is set to 100 ° C./s or more and 1000 ° C./s or less.
  • the quenching is preferably quenching by water quenching.
  • the secondary cooling stop temperature is 100 ° C or less. When the secondary cooling stop temperature is higher than 100 ° C, it will cause a decrease in the area ratio of martensite due to insufficient quenching of austenite during rapid cooling, and a decrease in material strength due to self-tempering of martensite generated by rapid cooling. It is not preferable.
  • a tempering process in which the residence time in the temperature range from 100 ° C to 300 ° C is 120 to 1800 seconds, and the tempering treatment is performed at a temperature of 100 ° C for the tempering of martensite. Re-heating to a temperature of 300 ° C or lower and tempering for 120 to 1800 seconds in the temperature range of 100 to 300 ° C.
  • This tempering softens martensite and improves workability.
  • tempering is performed at less than 100 ° C., the softening of martensite is insufficient, the effect of improving workability cannot be expected, and the hardness difference from ferrite becomes large, so that the stretch flangeability deteriorates.
  • the residence time is less than 120 seconds, the martensite is not sufficiently softened in the temperature range from 100 ° C. to 300 ° C., and therefore an effect of improving workability cannot be expected.
  • the residence time exceeds 1800 seconds, the strength is remarkably lowered due to excessive softening of martensite, and the manufacturing cost is increased due to an increase in reheating time, which is not preferable.
  • the Si oxide and Si-Mn oxide on the steel sheet surface are removed, and the chemical conversion processability is improved.
  • a non-oxidizing acid as the pickling solution, unlike the pickling solution used in pickling.
  • Pickling can be performed by a conventional method, and the conditions are not particularly limited. For example, any one of nitric acid, hydrochloric acid, hydrofluoric acid, sulfuric acid and an acid obtained by mixing two or more of them can be used.
  • the tempered steel sheet is pickled with a strong acid such as nitric acid with a concentration of more than 50 g / L and less than 200 g / L, for example.
  • non-oxidizing acid examples include hydrochloric acid, sulfuric acid, phosphoric acid, pyrophosphoric acid, formic acid, acetic acid, citric acid, hydrofluoric acid, oxalic acid, and acids obtained by mixing two or more of these.
  • hydrochloric acid having a concentration of 0.1 to 50 g / L sulfuric acid having a concentration of 0.1 to 150 g / L
  • an acid in which 0.1 to 20 g / L hydrochloric acid and 0.1 to 60 g / L sulfuric acid are mixed can be suitably used.
  • a high-strength cold-rolled steel sheet having a tensile strength (TS) of 1300 MPa or more and excellent in chemical conversion property and workability is produced. Since the high-strength cold-rolled steel sheet of the present invention is excellent in plate shape (flatness) after annealing, a process for correcting the shape of the steel sheet, such as rolling and leveler processing, is not necessarily required. From the viewpoint of adjusting the surface roughness, there is no problem even if the annealed steel sheet is rolled at an elongation of about several percent.
  • the high-strength cold-rolled steel sheet of the present invention does not affect the material depending on the plating treatment or the composition of the plating bath, so as a plating treatment, a hot dip galvanizing treatment, an alloyed hot dip galvanizing treatment, an electrogalvanizing treatment Any of the treatments can be applied.
  • Test steels A to R having the composition shown in Table 1 were vacuum-melted into slabs, and then hot rolled under the conditions shown in Table 2 to obtain hot-rolled steel sheets.
  • the hot-rolled steel sheet was pickled to remove the surface scale, and then cold-rolled (rolling rate: 60%).
  • continuous annealing and tempering treatment were performed under the conditions described in Table 2, and pickling and re- pickling were performed.
  • Ac 1 point and Ac 3 point were obtained from a transformation expansion curve obtained at an average heating rate of 3 ° C./s using a thermal expansion measuring device.
  • the two-phase volume ratio was determined by the point counting method based on the SEM image at a magnification of 1000 times, and the particle size of each phase was determined by the line segment method. The obtained volume ratio was defined as the area ratio.
  • the tensile test was carried out at a strain rate of 3.3 ⁇ 10 ⁇ 3 s ⁇ 1 by cutting a JIS No. 5 test piece (distance between gauge points: 50 mm, width of parallel part: 25 mm) parallel to the rolling direction. Total elongation was measured by test piece butt after fracture.
  • the hole expansion test was performed with a 100mm x 100mm size test piece, and after punching a circular hole of ⁇ 10mm (d 0 ), the apex angle was 60 ° with a 75mm inner diameter die pressed with a wrinkle holding force of 9tons.
  • the conical punch was pushed up from below against the hole, and the hole diameter (d) was measured when a plate thickness through crack occurred at the hole edge.
  • the hole expansion rate: (lambda) (%) defined by following Formula was calculated
  • the test was performed so that the hole punching and the hole expansion were in the same direction with the surface where burrs were generated by punching as the upper side (JIS 2256 compliant).
  • ⁇ (%) ⁇ (d-d0) / d0 ⁇
  • d0 initial hole diameter
  • d hole diameter at the time when the crack penetrates the plate thickness.
  • the steel sheet surface coverage of the oxide mainly composed of Si is mainly composed of Si in the same manner as above by observing 5 fields of view at 1000 times on the steel sheet surface using SEM and analyzing the same field of view by EDX.
  • the oxide was identified, and the coverage was determined by a point counting method (a method in which 15 straight lines were drawn in each of the vertical and horizontal directions of the SEM image to determine the probability that an Si-based oxide was present at the intersection (225 points)).
  • chemical conversion treatment was performed using a commercially available chemical treatment chemical (Nippon Parkerizing Co., Ltd., Palbond PB-L3065 (registered trademark)) at a bath temperature of 35 ° C and a treatment time of 120 seconds.
  • the processed steel sheet surface is observed with 5 fields using SEM at a magnification of 500 times, and when all the 5 fields have a uniform conversion crystal with an area ratio of 95% or more, the chemical conversion processability is good.
  • the chemical conversion treatment ability was evaluated as inferior “ ⁇ ” when a defect with an area ratio exceeding 5% was observed even in the visual field.
  • the examples that meet the conditions of the present invention have excellent strength and high tensile strength (TS) of 1300 MPa or more, elongation (EL) of 10% or more, and hole expansion ratio ( ⁇ ) of 30% or more. Workability is obtained. Moreover, it is excellent in chemical conversion treatment property.
  • No. 12 is a comparative example in which the C content is higher than the range of the present invention. Since the C content is high, the strength of martensite is increased and the balance between strength and ductility is excellent, but it can be seen that the stretch flangeability is extremely low due to the hardness difference between ferrite and martensite.
  • No. 13 and 14 are comparative examples in which the Si content is outside the scope of the present invention.
  • No. 13 does not satisfy the chemical conversion treatment property because Si oxide exists on the steel plate surface even after the two-step pickling treatment.
  • No. 14 does not have a predetermined elongation.
  • No. 15 and 16 are comparative examples in which the Mn content is outside the scope of the present invention. Since Mn is an element that greatly fluctuates the martensite fraction, No. 15 having a high content does not achieve a predetermined elongation. In No. 16 having a low content, the martensite fraction is small, so that a predetermined strength is not obtained.
  • Nos. 18 to 23 are comparative examples in which the manufacturing conditions are outside the scope of the present invention.
  • No. 18 is a comparative example in which the component composition and production conditions are outside the scope of the present invention. In addition to the fact that the predetermined elongation cannot be obtained, the stretch flangeability and chemical conversion treatment properties are inferior.
  • No. 19 has a high annealing temperature, so the prescribed strength and elongation cannot be obtained.
  • No. 20 to 22 do not have a sufficient martensite fraction, and the prescribed strength is not obtained.
  • No. 23 does not have a sufficient martensite fraction and has poor stretch flangeability.
  • No. 24 is an example in which the pickling treatment after annealing is omitted.
  • Si oxide is present on the surface of the steel sheet and does not satisfy the chemical conversion processability.

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Abstract

L'invention concerne une tôle d'acier laminée à froid à haute résistance, dont la résistance à la traction est supérieure ou égale à 1300 MPa, et présentant d'excellentes aptitudes au traitement de conversion chimique et au façonnage, ainsi que le procédé de fabrication de ladite tôle. La tôle d'acier contient de 0,15 % à 0,22 % de C, de 1,0 % à 2,0 % de Si, de 1,7 % à 2,5 % de Mn, 0,05 % ou moins de P, 0,02 % ou moins de S, de 0,01 % à 0,05 % d'Al, 0,005 % ou moins de N et 0,01 % ou moins de O. La tôle d'acier satisfait la relation [Si]/[Mn] ≥ 0,5 ([Si] est la teneur en Si et [Mn] la teneur en Mn), le reste contenant du fer et les inévitables impuretés. En pourcentage de surface, la structure contient de 60 % à moins de 100 % de martensite revenue et 5 % ou moins (0 % inclus) d'austénite résiduelle, le reste étant de la ferrite. La taille moyenne des grains cristallins de ferrite est inférieure ou égale à 3,5 µm. Le nombre d'oxydes composés de Si-Mn ayant un diamètre circulaire équivalent inférieur ou égal à 5 µm à la surface de la tôle d'acier est inférieur à 10/100 μm2. La couverture de la surface de la tôle d'acier en oxydes contenant principalement du Si est inférieure ou égale à 1 %.
PCT/JP2016/000778 2015-03-18 2016-02-16 Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication WO2016147549A1 (fr)

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JP2017506047A JP6210175B2 (ja) 2015-03-18 2016-02-16 高強度冷延鋼板およびその製造方法
CN201680016385.8A CN107429344A (zh) 2015-03-18 2016-02-16 高强度冷轧钢板及其制造方法
MX2017011825A MX2017011825A (es) 2015-03-18 2016-02-16 Lamina de acero laminada en frio de alta resistencia y metodo para producir la misma.
US15/556,448 US20180037969A1 (en) 2015-03-18 2016-02-16 High-strength cold-rolled steel sheet and method of producing the same
EP16764383.2A EP3272892B1 (fr) 2015-03-18 2016-02-16 Tôle d'acier laminée à froid à haute résistance et son procédé de fabrication
KR1020177025309A KR101998652B1 (ko) 2015-03-18 2016-02-16 고강도 냉연 강판 및 그의 제조 방법

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JP6388099B1 (ja) * 2017-12-15 2018-09-12 新日鐵住金株式会社 鋼板、溶融亜鉛めっき鋼板および合金化溶融亜鉛めっき鋼板
WO2019021695A1 (fr) * 2017-07-25 2019-01-31 Ｊｆｅスチール株式会社 Tôle d'acier haute résistance laminée à froid et son procédé de fabrication
EP3626849A4 (fr) * 2017-05-19 2020-05-06 JFE Steel Corporation Procédé de fabrication d'une tôle d'acier haute résistance à placage au zinc fondu
JP2020125535A (ja) * 2019-02-04 2020-08-20 Ｊｆｅスチール株式会社 冷延鋼板及びその製造方法
WO2022168167A1 (fr) * 2021-02-02 2022-08-11 日本製鉄株式会社 Tôle d'acier mince
WO2023096453A1 (fr) * 2021-11-29 2023-06-01 주식회사 포스코 Tôle d'acier laminée à froid ultra-haute résistance ayant un excellent allongement et son procédé de fabrication

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MX2018003405A (es) * 2015-09-25 2018-06-06 Nippon Steel & Sumitomo Metal Corp Lamina de acero.
EP3399064B1 (fr) * 2016-02-18 2021-07-14 JFE Steel Corporation Tôle en acier laminée à froid hautement résistante
KR102398869B1 (ko) * 2018-03-30 2022-05-16 제이에프이 스틸 가부시키가이샤 냉연 강판 및 그의 제조 방법
KR102519960B1 (ko) * 2018-11-09 2023-04-10 제이에프이 스틸 가부시키가이샤 지르코늄계 화성 처리용 냉연 강판 및 그의 제조 방법 그리고 지르코늄계 화성 처리 강판 및 그의 제조 방법
KR102488156B1 (ko) * 2019-01-09 2023-01-16 제이에프이 스틸 가부시키가이샤 고강도 냉연 강판 및 그 제조 방법
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KR101998652B1 (ko) 2019-07-10
MX2017011825A (es) 2017-12-07
EP3272892A1 (fr) 2018-01-24
US20180037969A1 (en) 2018-02-08
JP6210175B2 (ja) 2017-10-11
CN107429344A (zh) 2017-12-01

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