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WO2013111556A1 - High-strength hot-rolled steel sheet and method for producing same - Google Patents

High-strength hot-rolled steel sheet and method for producing same Download PDF

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Publication number
WO2013111556A1
WO2013111556A1 PCT/JP2013/000257 JP2013000257W WO2013111556A1 WO 2013111556 A1 WO2013111556 A1 WO 2013111556A1 JP 2013000257 W JP2013000257 W JP 2013000257W WO 2013111556 A1 WO2013111556 A1 WO 2013111556A1
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WO
WIPO (PCT)
Prior art keywords
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steel sheet
rolled steel
hot
strength
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PCT/JP2013/000257
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French (fr)
Japanese (ja)
Inventor
典晃 ▲高▼坂
船川　義正
重見　將人
英和 大久保
篤謙 金村
Original Assignee
Ｊｆｅスチール株式会社
Priority date (The priority date is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the date listed.)
Filing date
Publication date
Application filed by Ｊｆｅスチール株式会社 filed Critical Ｊｆｅスチール株式会社
Priority to CN201380005846.8A priority Critical patent/CN104053806B/en
Priority to KR1020147019785A priority patent/KR20140103340A/en
Priority to JP2013555188A priority patent/JP5565534B2/en
Priority to EP13740782.1A priority patent/EP2808413B1/en
Priority to US14/374,124 priority patent/US20150030880A1/en
Publication of WO2013111556A1 publication Critical patent/WO2013111556A1/en

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    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following 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
    • C21D8/0226—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
    • C21D8/0278—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips involving a particular surface treatment
    • C21D8/0284—Application of a separating or insulating coating
    • 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
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • 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
    • C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
    • 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
    • C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
    • 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
    • C22C38/007—Ferrous alloys, e.g. steel alloys containing silver
    • 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
    • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/10—Ferrous alloys, e.g. steel alloys containing cobalt
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
    • 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
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/54—Ferrous alloys, e.g. steel alloys containing chromium with nickel with boron
    • 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
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • 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
    • C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
    • C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
    • C23C2/02—Pretreatment of the material to be coated, e.g. for coating on selected surface areas
    • C23C2/022—Pretreatment of the material to be coated, e.g. for coating on selected surface areas by heating
    • C23C2/0224—Two or more thermal pretreatments
    • 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
    • C21D2211/004—Dispersions; Precipitations
    • 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
    • Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
    • Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
    • Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
    • Y10T428/00—Stock material or miscellaneous articles
    • Y10T428/12—All metal or with adjacent metals
    • Y10T428/12493—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, joint, etc.]
    • Y10T428/12771—Transition metal-base component
    • Y10T428/12785—Group IIB metal-base component
    • Y10T428/12792—Zn-base component
    • Y10T428/12799—Next to Fe-base component [e.g., galvanized]

Definitions

  • the present invention relates to a high-strength hot-rolled steel sheet having high tensile strength (TS): 980 MPa or more and excellent workability (particularly bending workability) useful for the use of automobile members and a method for producing the same.
  • TS tensile strength
  • Patent Document 1 proposes a technique for making a steel plate composition containing C: 0.05 to 0.20% by mass%, Nb: 0.1 to 1.0%, and having a solid solution C content of 0.03% or less. ing. And according to the technique proposed in Patent Document 1, by limiting the amount of dissolved C in the component system containing Nb and C, the matrix is a soft ferrite phase and the matrix is a hard second phase. It is said that a certain NbC-dispersed steel sheet structure is obtained, and a wear-resistant steel sheet having excellent bending workability is obtained.
  • the steel sheet composition is in mass%, C: 0.02 to 0.2%, Si: 0.01 to 1.0%, Mn: 0.1 to 2.0%, P: 0.2% or less, sol.Al: 0.001 to 0.5%, Ti: 0.1% or less, Nb: 0.1% or less, V: 0.5% or less, Mo: 0.5% or less, and Ti + Nb: 0.1% or less, steel sheet structure as ferrite main phase structure, and steel sheet surface
  • a technique for defining the average grain size of ferrite at a depth of 1/4 of the plate thickness and the rate of increase of the average grain size at 700 ° C. has been proposed. And according to the technique proposed by patent document 2, it is supposed that the steel plate excellent in workability will be obtained.
  • the technique proposed in Patent Document 1 substantially strengthens the steel sheet by dispersing NbC, and it is difficult to obtain a steel sheet having a tensile strength of 980 MPa or more with the technique using NbC. .
  • the amount of particle dispersion strengthening obtained by dispersing precipitates increases with an increase in the carbonized body volume fraction, but NbC increases the carbonized body volume fraction because its solubility product in steel is small and the atomic density is large. It is because it cannot be done.
  • Ti and V are added as precipitation strengthening elements to steel, but the contents of Ti and V forming carbides are small, or appropriate addition is not made. After all, the tensile strength of the steel sheet does not reach 980 MPa.
  • an object of the present invention is to provide a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability.
  • the present inventors focused on a technique for strengthening by dispersing fine carbides in a ferrite single-structure steel plate having good workability.
  • various factors affecting bending workability were studied.
  • the inventors of the present invention are extremely effective to disperse fine carbides in the ferrite phase in order to make the originally soft ferrite phase into a hard ferrite single-structure steel sheet.
  • Ti is the most suitable element.
  • the present inventors examined a means for imparting excellent bending workability while maintaining the steel plate strength of the high strength hot rolled steel plate having a tensile strength of 980 MPa or more to which Ti and V are added in combination as described above.
  • the surface properties of the steel sheet surface ⁇ appearance quality
  • the amount of solid solution elements and voids that are the starting points of voids that reduce the workability of the steel plate are reduced as much as possible. I found out that it was necessary.
  • the tensile strength is 980 MPa or more and It has been found that a hot-rolled steel sheet having excellent bending workability can be obtained.
  • the present invention has been completed based on the above findings, and the gist thereof is as follows.
  • the ferrite phase has an area ratio of 95% or more, the ferrite phase has an average crystal grain size of 8 ⁇ m or less, and the ferrite grain has an average grain size of less than 10 nm, and a tensile strength of 980 MPa.
  • the composition further contains at least one of Mo, W, Zr, and Hf. : 0.05% or less, Zr: 0.05% or less, Hf: High-strength hot-rolled steel sheet excellent in bendability characterized by being limited to 0.05% or less.
  • Heating the steel material subjecting it to hot rolling consisting of rough rolling and finish rolling, cooling after completion of finish rolling, winding, and hot rolling steel sheet,
  • the steel material in mass%, C: 0.06% to 0.1%, Si: 0.09% or less, Mn: 0.7% or more and 1.3% or less, P: 0.03% or less, S: 0.01% or less, Al: 0.1% or less, N: 0.01% or less, Ti: 0.14% to 0.20% V: 0.07% or more and 0.14% or less, with the balance being Fe and inevitable impurities,
  • the heating temperature of the steel material is 1100 ° C. or more and 1350 ° C. or less
  • the finish rolling temperature of the finish rolling is 850 ° C.
  • a method for producing a high-strength hot-rolled steel sheet having excellent bendability characterized by being 20 ° C / s or higher and a coiling temperature of the winding of 550 ° C or higher and 700 ° C or lower.
  • the composition further contains at least one of Mo, W, Zr, and Hf, and the content thereof is Mo: 0.05% or less, W : 0.05% or less, Zr: 0.05% or less, and Hf: 0.05% or less.
  • a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability suitable for the use of structural members of automobiles can be obtained.
  • the effect is remarkable such as making it possible, and further application development of the high-strength hot-rolled steel sheet becomes possible, and there is a remarkable industrial effect.
  • the hot-rolled steel sheet of the present invention has a structure in which the area ratio of the ferrite phase is 95% or more, the average crystal grain size of the ferrite phase is 8 ⁇ m or less, and the average grain size of carbides in the ferrite phase crystal grains is less than 10 nm. Have.
  • the matrix metal structure of the hot-rolled steel sheet is preferably a ferrite single-phase structure excellent in workability.
  • a second phase structure such as a bainite phase, martensite phase, cementite, or pearlite is mixed into the steel sheet structure, voids are generated at the interface between the ferrite phase and the second phase structure, which have different hardnesses, and the bending workability of the steel sheet is reduced.
  • it is intended to ensure the desired strength of the steel sheet by precipitating carbides such as Ti and V in the steel sheet, but most of these carbides are finished in the finish rolling in the hot-rolled steel sheet manufacturing process.
  • the metal structure of the hot-rolled steel sheet is a ferrite single phase structure, but if the ferrite area ratio is 95% or more even if it is not completely a ferrite single phase, the desired strength (Tensile strength: 980 MPa or more) is obtained, so the area ratio of the ferrite phase is 95% or more. Preferably it is 98% or more.
  • examples of the structure other than the ferrite phase that can be contained in the steel sheet include cementite, pearlite, bainite phase, and martensite phase.
  • the steel sheet properties bending workability, etc.
  • Average grain size of ferrite phase 8 ⁇ m or less
  • the average grain size of ferrite exceeds 8 ⁇ m, it tends to be a mixed grain size microstructure. In the mixed grain structure, stress tends to concentrate on coarse ferrite grains during bending, so that the bending workability of the steel sheet is significantly reduced. Therefore, the upper limit of the average crystal grain size of the ferrite phase is 8 ⁇ m. Preferably it is 6 micrometers or less, More preferably, it is 4.5 micrometers or less.
  • Carbides in ferrite crystal grains in the hot-rolled steel sheet of the present invention it is essential to finely precipitate carbides in the ferrite phase crystal grains from the viewpoint of ensuring strength.
  • Examples of the carbide finely precipitated in the ferrite phase grains in the present invention include Ti carbide, V carbide, and Ti and V composite carbide, or those containing Nb, W, Mo, Hf, and Zr in the carbide. .
  • Most of these carbides are carbides that precipitate at the interface at the same time as the transformation from austenite to ferrite in the cooling process after finish rolling in the hot-rolled steel sheet manufacturing process.
  • Average particle diameter of carbides in ferrite crystal grains less than 10 nm
  • the aforementioned carbides mainly composite carbides of Ti and V
  • the carbides are fine.
  • the average particle size of the carbide dispersed in the ferrite crystal grains is set to less than 10 nm.
  • it is less than 7 nm, More preferably, it is 5 nm or less.
  • C 0.06% or more and 0.1% or less C is combined with Ti, V, or Nb and is finely distributed in the steel sheet as carbides. That is, C is an element that forms fine carbides and remarkably strengthens the ferrite structure, and is an essential element for strengthening the hot-rolled steel sheet. In order to obtain a high-strength steel sheet having a tensile strength of 980 MPa or more, the C content needs to be at least 0.06%.
  • the C content exceeds 0.1%, a large amount of cementite precipitates and the bending workability of the steel sheet is lowered. This is because microvoids are easily generated at the cementite / matrix (ferrite) interface, and these microvoids cause cracks in the bent portion of the steel sheet. Therefore, the C content is set to 0.06% or more and 0.1% or less. Preferably they are 0.07% or more and 0.09% or less.
  • Si 0.09% or less Si is positively contained in conventional high-strength steel sheets as an effective element for improving the steel sheet strength without reducing ductility (elongation).
  • Si is easy to concentrate on the steel sheet surface, and forms firelite (Fe 2 SiO 4 ) on the steel sheet surface. Since this firelite is formed in a wedge shape on the surface of the steel sheet, it becomes a starting point of cracking when the steel sheet is bent. Therefore, in the present invention, it is desirable to reduce the Si content as much as possible, but 0.09% is acceptable, so the upper limit of the Si content is 0.09%. Preferably it is 0.06% or less. Note that the Si content may be reduced to the impurity level.
  • Mn 0.7% or more and 1.3% or less Mn is an element effective for increasing the strength of a hot-rolled steel sheet in order to refine the carbides precipitated in the crystal grains of the ferrite phase.
  • Mn 0.7% or more and 1.3% or less Mn is an element effective for increasing the strength of a hot-rolled steel sheet in order to refine the carbides precipitated in the crystal grains of the ferrite phase.
  • Mn has the effect of lowering the transformation temperature from austenite to ferrite of steel, so by containing a predetermined amount of Mn, the transformation temperature is lowered to the coiling temperature range described later, Carbide can be deposited simultaneously with the winding of the steel sheet. And the carbide
  • the Mn content needs to be at least 0.7%.
  • the Mn content exceeds 1.3%, the workability of the steel sheet due to solute Mn is significantly reduced, so that the desired bending workability cannot be obtained. Therefore, the Mn content is 0.7% or more and 1.3% or less. Preferably they are 0.8% or more and 1.2% or less.
  • P 0.03% or less
  • P is a harmful element that segregates at the grain boundary and becomes the starting point of grain boundary cracking during processing, and degrades the bending workability of the steel sheet, so it is preferably reduced as much as possible. Therefore, in the present invention, the P content is set to 0.03% or less in order to avoid the above problems. Preferably it is 0.02% or less. The P content may be reduced to the impurity level.
  • S 0.01% or less S is present as an inclusion such as MnS in steel. Since the inclusions are hard, the interface between the matrix and the inclusions becomes a starting point of voids during bending of the steel sheet, thereby reducing the bending workability of the steel sheet. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, and set it to 0.01% or less. Preferably it is 0.008% or less. There is no problem even if the S content is zero.
  • Al 0.1% or less
  • Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is desirable to contain 0.02% or more. However, if the Al content exceeds 0.1%, an adverse effect on bending workability due to inclusions such as alumina becomes obvious. Therefore, the Al content is 0.1% or less. Preferably it is 0.08% or less.
  • N 0.01% or less N is combined with Ti, which is a carbide-forming element, at the steelmaking stage to form coarse Ti nitrides, which impede the formation of fine carbides and significantly reduce the steel sheet strength. Furthermore, voids are likely to occur at the interface between the matrix and coarse Ti nitride during bending of the steel sheet, which adversely affects the bending workability of the steel sheet. Therefore, the N content is preferably reduced as much as possible, and is 0.01% or less. Preferably it is 0.008% or less. There is no problem even if the N content is zero.
  • Ti 0.14% or more and 0.20% or less Ti is an element that contributes to increasing the strength of steel sheets by forming carbides with C.
  • the Ti content needs to be 0.14% or more.
  • coarse Ti carbide cannot be dissolved by heating the steel material (slab) before hot rolling, and finally obtained ( Coarse Ti carbide remains on the hot-rolled steel sheet after winding.
  • the Ti content is 0.14% or more and 0.20% or less. Preferably it is 0.15% or more and 0.19% or less.
  • V 0.07% or more and 0.14% or less
  • V is an element that contributes to increasing the strength of steel sheets by forming carbides with C.
  • V is effective for increasing the strength of the steel sheet because it combines with Ti to form a fine composite carbide.
  • the V content needs to be 0.07% or more.
  • V content when the V content is higher than the Ti content, V becomes difficult to precipitate and the amount of V remaining in the steel sheet as a solid solution state increases. Since V in a solid solution state deteriorates the bending workability of the steel sheet, the V content needs to be Ti content or less, that is, 0.14% or less. Therefore, the V content is 0.07% or more and 0.14% or less. Preferably they are 0.08% or more and 0.13% or less.
  • Nb 0.01% or more and 0.05% or less
  • the above is the basic component in the present invention.
  • Nb 0.01% or more and 0.05% or less may be further contained.
  • Nb has the effect of inhibiting the recrystallization of austenite grains before transformation from austenite to ferrite in the hot rolling process when producing a hot rolled steel sheet having a ferrite single-phase structure substantially to make an unrecrystallized structure There is. In the non-recrystallized structure, strain energy due to hot rolling is easily accumulated, and the number of nucleation sites of the ferrite phase increases.
  • the Nb content is preferably set to 0.01% or more.
  • the Nb content is preferably 0.05% or less. More preferably, it is 0.02% or more and 0.04% or less.
  • Mo 0.05% or less
  • W 0.05% or less
  • Zr 0.05% or less
  • Hf 0.05% or less
  • Mo, W, Zr, and Hf are elements that form carbides and contribute to increasing the strength of the steel sheet, but a large amount remains as a solid solution. These solid solution elements deteriorate the workability of the matrix and adversely affect the bending workability of the steel sheet.
  • Mo, W, Zr, and Hf have a low rate of precipitation with respect to the content, and a large amount remains in the steel sheet as a solid solution element. For this reason, it is desirable to reduce these contents as much as possible, but 0.05% is acceptable, so the upper limit was set to 0.05%. Preferably it is 0.03% or less. Note that the contents of Mo, W, Zr, and Hf may be zero.
  • a plating layer By forming a plating layer on the surface, the corrosion resistance of the hot-rolled steel sheet is improved, and it becomes possible to apply it to automobile parts used in severe corrosive environments.
  • the type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied.
  • the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer and the like can be mentioned as suitable examples, but of course, it is not limited to these, and any conventionally known one can be applied. is there.
  • a steel material (steel slab) having the above composition is heated, subjected to hot rolling consisting of rough rolling and finish rolling, cooled after completion of finish rolling, and wound into a hot rolled steel sheet. Then, the heating temperature of the steel material is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, and the cooling is started within 3 seconds after finishing the finish rolling, and the average cooling of the cooling The speed is 20 ° C./s or more, and the winding temperature of the winding is 550 ° C. or more and 700 ° C. or less.
  • 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. .
  • Heating temperature of steel material 1100 ° C or higher and 1350 ° C or lower
  • the steel material (steel slab) obtained as described above is subjected to rough rolling and finish rolling.
  • the steel material is heated prior to rough rolling. It is necessary to form a substantially homogeneous austenite phase and dissolve coarse carbides.
  • the heating temperature of the steel material is below 1100 ° C, coarse carbides do not dissolve, so the amount of carbides that are finely dispersed in the cooling and winding process after hot rolling is reduced, and the hot rolling finally obtained is reduced.
  • the strength of the steel sheet is significantly reduced.
  • the heating temperature exceeds 1350 ° C., the scale bites and deteriorates the surface appearance quality of the steel sheet.
  • the heating temperature of the steel material is set to 1100 ° C to 1350 ° C. Preferably they are 1150 degreeC or more and 1320 degrees C or less.
  • the steel material is heated. Direct rolling may be performed without any problem.
  • the rough rolling conditions are not particularly limited.
  • Finish rolling temperature 850 ° C or more
  • the finish rolling temperature is 850 ° C. or higher.
  • it is 870 degreeC or more.
  • the upper limit of the finish rolling temperature is not particularly defined, but the finish rolling temperature is determined by the slab reheating temperature, the sheet feed speed, and the steel plate thickness. Therefore, the upper limit of the finish rolling temperature is substantially 990 ° C. or less.
  • carbides are generated by strain-induced precipitation because the strain energy accumulated in the austenite phase is large. Since this carbide precipitates at a high temperature and is easy to coarsen, it is difficult to obtain a fine precipitate when strain-induced precipitation occurs. Therefore, in the present invention, for the purpose of suppressing strain-induced precipitation, it is necessary to start forced cooling immediately after the end of hot rolling, and cooling is started within 3 seconds at the latest after finishing rolling. Preferably it is within 2 seconds.
  • Average cooling rate 20 ° C./s or more
  • the longer the time that the steel sheet after finish rolling is maintained at a high temperature the more easily the coarsening of the carbide by strain-induced precipitation proceeds.
  • the austenite ⁇ ferrite transformation is suppressed by containing a predetermined amount of Mn in the steel sheet, but if the cooling rate is low, the ferrite transformation starts at a high temperature, and the carbide tends to become coarse. Therefore, it is necessary to rapidly cool after finish rolling, and in order to avoid the above problem, it is necessary to cool at an average cooling rate of 20 ° C./s or more.
  • a preferable average cooling rate is 40 ° C./s or more.
  • the cooling rate after finishing rolling is excessively increased, it is difficult to control the coiling temperature and it is difficult to obtain stable hot-rolled steel sheet strength. preferable.
  • Winding temperature 550 ° C. or higher and 700 ° C. or lower If the winding temperature is lower than 550 ° C., a sufficient amount of carbide cannot be obtained, and the steel sheet strength decreases. On the other hand, when the coiling temperature exceeds 700 ° C., the precipitated carbide is coarsened, so that the steel sheet strength is lowered. Accordingly, the coiling temperature range is 550 ° C. or higher and 700 ° C. or lower. Preferably they are 580 degreeC or more and 680 degrees C or less.
  • the hot-rolled steel sheet after being rolled by hot rolling does not change its properties even if the scale is attached to the surface or the scale is removed by pickling.
  • the above-described excellent characteristics are exhibited in any state.
  • the hot-rolled steel sheet after winding may be plated to form a plating layer on the surface of the hot-rolled steel sheet.
  • the hot-rolled steel sheet of the present invention has little material fluctuation even when subjected to heat treatment up to 740 ° C. for a short time. Therefore, for the purpose of imparting corrosion resistance to the steel sheet, the hot-rolled steel sheet of the present invention can be plated, and a plating layer can be provided on the surface thereof. Since it can be manufactured even at a heating temperature of 740 ° C. or lower in the plating process, the effect of the present invention described above is not impaired even if the hot-rolled steel sheet of the present invention is plated.
  • the type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied.
  • the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer and the like can be mentioned as suitable examples, but of course, it is not limited to these, and any conventionally known one can be applied. is there.
  • the plating method is not particularly limited, and any conventionally known method can be applied.
  • a method of immersing and pulling up the steel plate in a plating bath may be used.
  • the steel plate surface may be heated in a furnace such as a gas furnace to perform the alloying treatment.
  • the ferrite phase area ratio is 95% or more, the average grain size of the ferrite phase is 8 ⁇ m or less, the ferrite phase of the ferrite phase A hot-rolled steel sheet having a structure in which the carbide average particle diameter in the crystal grains is less than 10 nm is obtained.
  • the solid solution elements and coarse inclusions present in the steel sheet are reduced, and the strength is increased, so that the high strength hot rolling excellent in bending workability is achieved. It can be a steel plate.
  • the content of carbide forming elements (Ti and V, or even Nb, W, Mo, Hf, Zr) contained in the steel sheet is optimized, and the production conditions for the hot rolled steel sheet are specified. .
  • the above-mentioned carbide having an average particle diameter of less than 10 nm can be sufficiently precipitated in the ferrite crystal grains, and the tensile strength of the hot-rolled steel sheet is increased to 980 MPa or more while maintaining excellent bending workability. be able to.
  • the present invention is preferably applied to a hot rolled steel sheet having a tensile strength of 1100 MPa or less, and more preferably applied to a hot rolled steel sheet having a tensile strength of 1052 MPa or less.
  • a steel material (steel slab) having a composition shown in Table 1 and having a thickness of 250 mm was hot rolled under the hot rolling conditions shown in Table 2 to obtain a hot rolled steel sheet having a thickness of 1.4 to 3.2 mm.
  • the cooling rate described in Table 2 is an average cooling rate from the finish rolling temperature to the winding temperature.
  • a part of the obtained hot-rolled steel sheet is passed through a hot dip galvanizing line with an annealing temperature of 720 ° C, and then immersed in a 460 ° C plating bath (plating composition: Zn-0.13 mass% Al).
  • a hot-dip galvanized material (GI material) was used.
  • Samples are taken from the hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material) obtained as described above, and subjected to structure observation, tensile test, and bending test.
  • the area ratio of the ferrite phase and the type of structure other than the ferrite phase The area ratio, the average crystal grain size of the ferrite phase, the average grain size of the carbide, the yield strength, the tensile strength, the elongation, and the critical bending radius were obtained.
  • the test method was as follows.
  • the area ratio of the ferrite phase was evaluated by the following method. The central portion of the plate thickness in the cross section parallel to the rolling direction was photographed for 10 fields of view by corroding the appearance of corrosion by 5% nital 400 times with a scanning optical microscope.
  • the ferrite phase is a structure having a form in which corrosion marks and cementite are not observed in the grains. Further, using polygonal ferrite, bainitic ferrite, acicular ferrite, and granular ferrite as ferrite, the area ratio of the ferrite phase, the average grain size of the ferrite phase, and the average grain size of the carbide in the crystal grains of the ferrite phase were derived.
  • the area ratio of the ferrite phase was obtained by separating the ferrite phase from the ferrite phase other than the ferrite phase such as bainite and martensite by image analysis, and obtaining the area ratio of the ferrite phase with respect to the observation field. At this time, the grain boundary observed as a linear form was counted as a part of the ferrite phase.
  • Table 3 shows the area ratio of the obtained ferrite phase.
  • the average crystal grain size of the ferrite phase was obtained by enlarging the above 400 times and taking three representative photographs by drawing three horizontal lines and three vertical lines, respectively, by a cutting method in accordance with ASTM E 112-10. Was the final average crystal grain size. Table 3 shows the average crystal grain size obtained.
  • the average particle size of the carbides in the ferrite phase grains was measured using a transmission electron microscope (magnification: 135000 times) by preparing a sample from the center of the thickness of the obtained hot rolled steel sheet using a thin film method. It calculated
  • All examples of the present invention are hot-rolled steel sheets having a tensile strength of TS: 980 MPa or more, excellent bending workability, and both strength and workability. On the other hand, it was found that the comparative example out of the scope of the present invention did not ensure a predetermined high strength or did not obtain good bending workability.
  • a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability suitable for the use of structural members of automobiles can be obtained. It becomes possible to achieve both.

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Abstract

The present invention addresses the problem of providing: a high-strength hot-rolled steel sheet which has excellent strength and excellent processability (in particular, bending processability) at the same time; and a method for producing the high-strength hot-rolled steel sheet. In order to solve the above-mentioned problem, the present invention is characterized by: having a specific composition; having an area ratio of a ferrite phase of 95% or more; having an average crystal grain diameter of the ferrite phase of 8 μm or less; having a structure wherein the average particle diameter of carbides within the crystal grains of the ferrite phase is less than 10 nm; and having a tensile strength of 980 MPa or more.

Description

高強度熱延鋼板及びその製造方法High strength hot-rolled steel sheet and manufacturing method thereof
　本発明は、自動車用部材の使途に有用な、引張強さ(TS)：980MPa以上の高い強度と優れた加工性（特に曲げ加工性）を兼ね備えた高強度熱延鋼板及びその製造方法に関する。 The present invention relates to a high-strength hot-rolled steel sheet having high tensile strength (TS): 980 MPa or more and excellent workability (particularly bending workability) useful for the use of automobile members and a method for producing the same.
　近年地球環境保全の観点から、CO2排出量の規制を目的として自動車業界全体で自動車の燃費改善が指向されている。自動車の燃費改善には、使用部材の薄肉化による自動車の軽量化が最も有効であるため、近年、自動車部品用素材としての高強度熱延鋼板の使用量が増加しつつある。一方、鋼板を素材とする自動車部材の多くは、プレス加工等によって成形されるため、自動車部品用鋼板には高強度に加えて優れた曲げ性（曲げ加工性）を有することも要求される。 In recent years, from the viewpoint of global environmental conservation, the automobile industry as a whole has been aimed at improving the fuel efficiency of automobiles for the purpose of regulating CO 2 emissions. In order to improve the fuel consumption of automobiles, it is most effective to reduce the weight of automobiles by thinning the members used. In recent years, the amount of high-strength hot-rolled steel sheets used as materials for automobile parts is increasing. On the other hand, since many automobile members made of steel plates are formed by press working or the like, steel plates for automobile parts are required to have excellent bendability (bendability) in addition to high strength.
　ただし、一般的に鉄鋼材料は高強度化に伴い延性が低下して加工性が劣化することから、引張強さを980MPa以上にまで高強度化した鋼板は、所望の部品形状に成形加工する際、様々な支障をきたすおそれがある。例えば、引張強さ：980MPa以上の鋼板にプレス加工を施す場合、曲げ加工部での割れやネッキング等の発生が顕著となるため、部品の成形が困難となる。
　そのような理由から、高強度鋼板を自動車部品等に適用する上で、所望の強度と優れた曲げ加工性を兼ね備えた高強度鋼板の開発が望まれており、現在までに様々な技術が提案されている。 However, since steel materials generally have a lower ductility and higher workability due to higher strength, steel sheets with higher tensile strength up to 980 MPa or more must be formed into a desired part shape. There is a risk of causing various troubles. For example, when a steel sheet having a tensile strength of 980 MPa or more is subjected to press working, cracks, necking, and the like occur in the bent portion, making it difficult to form the part.
For this reason, the development of high-strength steel sheets that have both the desired strength and excellent bending workability is desired when applying high-strength steel sheets to automobile parts, etc., and various technologies have been proposed to date. Has been.
　例えば、特許文献１では、鋼板組成を、質量％でC：0.05～0.20％、Nb：0.1～1.0％を含有し、かつ、固溶C量が0.03％以下である組成とする技術が提案されている。そして、特許文献１で提案された技術によると、NbとCを含有する成分系で固溶C量を制限することにより、マトリクスが軟質なフェライト相であり且つ該マトリクス中に硬質第二相であるNbCが分散した鋼板組織となり、優れた曲げ加工性を有する耐摩耗鋼板が得られるとされている。 For example, Patent Document 1 proposes a technique for making a steel plate composition containing C: 0.05 to 0.20% by mass%, Nb: 0.1 to 1.0%, and having a solid solution C content of 0.03% or less. ing. And according to the technique proposed in Patent Document 1, by limiting the amount of dissolved C in the component system containing Nb and C, the matrix is a soft ferrite phase and the matrix is a hard second phase. It is said that a certain NbC-dispersed steel sheet structure is obtained, and a wear-resistant steel sheet having excellent bending workability is obtained.
　また、特許文献２では、鋼板組成を質量％で、C：0.02～0.2％、Si：0.01～1.0％、Mn：0.1～2.0％、P：0.2％以下、sol.Al：0.001～0.5％、Ti：0.1％以下、Nb：0.1％以下、V：0.5％以下、Mo：0.5％以下、かつ、Ti＋Nb：0.1％以下を含有する組成とし、鋼板組織をフェライト主相組織とし、さらに、鋼板表面から板厚の1/4の深さにおけるフェライトの平均結晶粒径と、該平均結晶粒径の700℃における増加速度とを規定する技術が提案されている。そして、特許文献２で提案された技術によると、加工性に優れた鋼板が得られるとされている。 In Patent Document 2, the steel sheet composition is in mass%, C: 0.02 to 0.2%, Si: 0.01 to 1.0%, Mn: 0.1 to 2.0%, P: 0.2% or less, sol.Al: 0.001 to 0.5%, Ti: 0.1% or less, Nb: 0.1% or less, V: 0.5% or less, Mo: 0.5% or less, and Ti + Nb: 0.1% or less, steel sheet structure as ferrite main phase structure, and steel sheet surface A technique for defining the average grain size of ferrite at a depth of 1/4 of the plate thickness and the rate of increase of the average grain size at 700 ° C. has been proposed. And according to the technique proposed by patent document 2, it is supposed that the steel plate excellent in workability will be obtained.
特開２００７－２６２４２９号公報JP 2007-262429 A 特開２００８－１８９９７８号公報JP 2008-189978 A
　しかしながら、特許文献１で提案された技術は、NbCを分散させることで実質的に鋼板を強化するものであり、NbCを活用した当該技術では引張強さ980MPa以上の鋼板を得ることが困難である。析出物を分散させることで得られる粒子分散強化量は炭化物体積分率の増大に伴い上昇するが、NbCは鋼中への溶解度積が小さく原子密度が大きいことから、炭化物体積分率を大きくすることができないためである。
　また、特許文献２で提案された技術では、鋼に析出強化元素としてTiやVを添加しているが、炭化物を形成するTiとVの含有量が少ない、若しくは適正な添加がなされていないため、やはり鋼板の引張強さは980MPaに達していない。 However, the technique proposed in Patent Document 1 substantially strengthens the steel sheet by dispersing NbC, and it is difficult to obtain a steel sheet having a tensile strength of 980 MPa or more with the technique using NbC. . The amount of particle dispersion strengthening obtained by dispersing precipitates increases with an increase in the carbonized body volume fraction, but NbC increases the carbonized body volume fraction because its solubility product in steel is small and the atomic density is large. It is because it cannot be done.
Further, in the technique proposed in Patent Document 2, Ti and V are added as precipitation strengthening elements to steel, but the contents of Ti and V forming carbides are small, or appropriate addition is not made. After all, the tensile strength of the steel sheet does not reach 980 MPa.
　以上のように、従来技術では、引張強さが980MPa以上である高強度鋼板を得ることが困難であった。さらに、このような鋼板強度を維持しつつ優れた曲げ加工性を付与することはできなかった。
　本発明は、かかる事情に鑑みてなされたものであって、980MPa以上の引張強さを有し、曲げ加工性にも優れた高強度熱延鋼板を提供することを目的とする。 As described above, it has been difficult to obtain a high-strength steel sheet having a tensile strength of 980 MPa or more with the prior art. Furthermore, excellent bending workability could not be imparted while maintaining such steel plate strength.
The present invention has been made in view of such circumstances, and an object of the present invention is to provide a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability.
　上記課題を解決すべく、本発明者らは、加工性が良好なフェライト単一組織鋼板に微細な炭化物を分散させることで強化を図る技術に着目し、該鋼板の高強度化と加工性、特に曲げ加工性に及ぼす各種要因について鋭意検討した。　
　そして、本発明者らは、本来では軟質なフェライト相を、硬質なフェライト単一組織鋼板とするためには、フェライト相に微細な炭化物を分散させることが極めて有効であり、微細な炭化物を多く析出できる元素を模索した結果、Tiが最も適した元素であることを知見した。 In order to solve the above-mentioned problems, the present inventors focused on a technique for strengthening by dispersing fine carbides in a ferrite single-structure steel plate having good workability. In particular, various factors affecting bending workability were studied.
The inventors of the present invention are extremely effective to disperse fine carbides in the ferrite phase in order to make the originally soft ferrite phase into a hard ferrite single-structure steel sheet. As a result of searching for elements that can be precipitated, it was found that Ti is the most suitable element.
　しかしながら、Ti炭化物のみでは、フェライト単一組織である熱延鋼板の引張強さを980MPa以上とすることが困難であったことから、本発明者らは、Ti炭化物による分散析出強化を補強する手段について検討した。
　その結果、補強手段としてVを添加することに想到した。Vは鋼中での溶解度が大きいため、単独添加では析出し難いが、Ti炭化物と結合することで析出し易い状態となる。その結果、熱延鋼板の鋼素材に適正量のTi及びVを複合添加すると、Ti又はVを単独で添加する場合に比べて鋼板強度が飛躍的に高くなり、引張強さ980MPa以上の熱延鋼板が得られる。
　また、Vは、Ti炭化物と結合することで析出し易い状態となることから、Ti含有量をV含有量以上とすることも重要であることが明らかとなった。 However, with Ti carbide alone, it was difficult to make the tensile strength of the hot rolled steel sheet, which is a ferrite single structure, 980 MPa or more. Was examined.
As a result, they came up with adding V as a reinforcing means. V has a high solubility in steel, so it is difficult to precipitate when added alone, but it is likely to precipitate when combined with Ti carbide. As a result, when a proper amount of Ti and V is added to the steel material of the hot-rolled steel sheet, the steel sheet strength is dramatically higher than when Ti or V is added alone, and the hot-rolled steel with a tensile strength of 980 MPa or more is greatly increased. A steel plate is obtained.
Moreover, since V becomes a state which precipitates easily by couple | bonding with Ti carbide, it became clear that it is important to make Ti content more than V content.
　また、本発明者らは、上記の如くTi及びVを複合添加した引張強さ980MPa以上の高強度熱延鋼板について、鋼板強度を維持しつつ優れた曲げ加工性を付与する手段について検討した。その結果、曲げ加工性を付与するには、鋼板の表面性状（surface appearance quality）を改善し、さらに鋼板の加工性を低下させる固溶元素量やボイドの発生起点となる介在物を極力低減する必要があることを知見した。そして、さらに検討を進めた結果、Si含有量を極力低減したうえで、C、Mn、Ti及びV含有量の最適化を図った成分組成とすることで、引張強さが980MPa以上であり且つ優れた曲げ加工性を有する熱延鋼板が得られることを見出した。 Further, the present inventors examined a means for imparting excellent bending workability while maintaining the steel plate strength of the high strength hot rolled steel plate having a tensile strength of 980 MPa or more to which Ti and V are added in combination as described above. As a result, in order to impart bending workability, the surface properties of the steel sheet (surface 鋼板 appearance quality) are improved, and the amount of solid solution elements and voids that are the starting points of voids that reduce the workability of the steel plate are reduced as much as possible. I found out that it was necessary. And, as a result of further investigation, after reducing the Si content as much as possible, by making the component composition that optimizes the C, Mn, Ti and V content, the tensile strength is 980 MPa or more and It has been found that a hot-rolled steel sheet having excellent bending workability can be obtained.
　本発明は上記の知見に基づき完成されたものであり、その要旨は次のとおりである。
［１］質量％で、
C：0.06％以上0.1％以下、
Si：0.09％以下、
Mn：0.7％以上1.3％以下、
P：0.03％以下、
S：0.01％以下、
Al：0.1％以下、
N：0.01％以下、
Ti：0.14％以上0.20％以下
V：0.07％以上0.14％以下
を含有し、残部がFe及び不可避的不純物である組成からなり、
　フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が8μm以下、該フェライト相の結晶粒内の炭化物の平均粒子径が10nm未満である組織を有し、引張強さが980MPa以上であることを特徴とする、曲げ性に優れた高強度熱延鋼板。 The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%
C: 0.06% to 0.1%,
Si: 0.09% or less,
Mn: 0.7% or more and 1.3% or less,
P: 0.03% or less,
S: 0.01% or less,
Al: 0.1% or less,
N: 0.01% or less,
Ti: 0.14% to 0.20%
V: Containing 0.07% or more and 0.14% or less, with the balance being Fe and inevitable impurities,
The ferrite phase has an area ratio of 95% or more, the ferrite phase has an average crystal grain size of 8 μm or less, and the ferrite grain has an average grain size of less than 10 nm, and a tensile strength of 980 MPa. A high-strength hot-rolled steel sheet excellent in bendability characterized by the above.
［２］前記［１］において、前記組成に加えてさらに、質量％でNb：0.01％以上0.05％以下を含有することを特徴とする曲げ性に優れた高強度熱延鋼板。 [2] A high-strength hot-rolled steel sheet excellent in bendability according to [1], further containing Nb: 0.01% or more and 0.05% or less by mass% in addition to the composition.
［３］前記［１］又は［２］において、前記組成に加えてさらに、Mo、W、Zr、Hfのいずれか１種以上を含有し、これらの含有量を、Mo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下に制限したことを特徴とする曲げ性に優れた高強度熱延鋼板。 [3] In the above [1] or [2], in addition to the above composition, the composition further contains at least one of Mo, W, Zr, and Hf. : 0.05% or less, Zr: 0.05% or less, Hf: High-strength hot-rolled steel sheet excellent in bendability characterized by being limited to 0.05% or less.
［４］前記［１］～［３］のいずれかにおいて、前記組成に加えてさらに、質量％で、O（酸素）、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、B、Ni、Cr、Sb、Cu、Sn、Mg及びCaのうちの１種以上を合計で0.2％以下含有することを特徴とする曲げ性に優れた高強度熱延鋼板。 [4] In any one of the above [1] to [3], in addition to the composition, O (oxygen), Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, Pd, Pt, Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, B, Ni, Cr, Sb, Cu, Sn, A high-strength hot-rolled steel sheet excellent in bendability characterized by containing one or more of Mg and Ca in a total amount of 0.2% or less.
［５］前記［１］～［４］のいずれかにおいて、前記鋼板表面に、めっき層をさらに有することを特徴とする曲げ性に優れた高強度熱延鋼板。 [5] The high-strength hot-rolled steel sheet with excellent bendability according to any one of [1] to [4], further comprising a plating layer on the steel sheet surface.
［６］前記［５］において、前記めっき層が亜鉛めっき層であることを特徴とする曲げ性に優れた高強度熱延鋼板。 [6] A high-strength hot-rolled steel sheet excellent in bendability according to [5], wherein the plated layer is a galvanized layer.
［７］前記［５］において、前記めっき層が合金化亜鉛めっき層であることを特徴とする曲げ性に優れた高強度熱延鋼板。 [7] A high-strength hot-rolled steel sheet excellent in bendability according to [5], wherein the plated layer is an alloyed galvanized layer.
［８］鋼素材を加熱し、粗圧延と仕上圧延（finish rolling）からなる熱間圧延を施し、仕上圧延終了後、冷却し、巻き取り、熱延鋼板とするにあたり、
　前記鋼素材を、質量％で、
C：0.06％以上0.1％以下、
Si：0.09％以下、
Mn：0.7％以上1.3％以下、
P：0.03％以下、
S：0.01％以下、
Al：0.1％以下、
N：0.01％以下、
Ti：0.14％以上0.20％以下
V：0.07％以上0.14％以下
を含有し、残部がFe及び不可避的不純物からなる組成とし、
　前記鋼素材の加熱温度を1100℃以上1350℃以下とし、前記仕上圧延の仕上圧延温度を850℃以上とし、前記冷却を仕上圧延終了後から3秒以内に開始し、前記冷却の平均冷却速度を20℃/s以上とし、前記巻き取りの巻取り温度（coiling temperature）を550℃以上700℃以下とすることを特徴とする、曲げ性に優れた高強度熱延鋼板の製造方法。 [8] Heating the steel material, subjecting it to hot rolling consisting of rough rolling and finish rolling, cooling after completion of finish rolling, winding, and hot rolling steel sheet,
The steel material in mass%,
C: 0.06% to 0.1%,
Si: 0.09% or less,
Mn: 0.7% or more and 1.3% or less,
P: 0.03% or less,
S: 0.01% or less,
Al: 0.1% or less,
N: 0.01% or less,
Ti: 0.14% to 0.20%
V: 0.07% or more and 0.14% or less, with the balance being Fe and inevitable impurities,
The heating temperature of the steel material is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling is set. A method for producing a high-strength hot-rolled steel sheet having excellent bendability, characterized by being 20 ° C / s or higher and a coiling temperature of the winding of 550 ° C or higher and 700 ° C or lower.
［９］前記［８］において、前記組成に加えてさらに、質量％でNb：0.01％以上0.05％以下を含有することを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [9] A method for producing a high-strength hot-rolled steel sheet excellent in bendability according to [8], further comprising Nb: 0.01% to 0.05% by mass% in addition to the composition.
［１０］前記［８］又は［９］において、前記組成に加えてさらに、Mo、W、Zr、Hfのいずれか１種以上を含有し、これらの含有量を、Mo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下に制限したことを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [10] In the above [8] or [9], in addition to the above composition, the composition further contains at least one of Mo, W, Zr, and Hf, and the content thereof is Mo: 0.05% or less, W : 0.05% or less, Zr: 0.05% or less, and Hf: 0.05% or less. A method for producing a high-strength hot-rolled steel sheet with excellent bendability.
［１１］前記［８］～［１０］のいずれかにおいて、前記組成に加えてさらに、質量％で、O（酸素）、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、B、Ni、Cr、Sb、Cu、Sn、Mg及びCaのうちの１種以上を合計で0.2％以下含有することを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [11] In any one of the above [8] to [10], in addition to the composition, O (oxygen), Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, Pd, Pt, Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, B, Ni, Cr, Sb, Cu, Sn, A method for producing a high-strength hot-rolled steel sheet excellent in bendability, comprising one or more of Mg and Ca in a total amount of 0.2% or less.
［１２］前記［８］～［１１］のいずれかにおいて、前記熱延鋼板の表面にめっき層を形成することを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [12] The method for producing a high-strength hot-rolled steel sheet excellent in bendability according to any one of [8] to [11], wherein a plating layer is formed on the surface of the hot-rolled steel sheet.
［１３］前記［１２］において、前記めっき層が亜鉛めっき層であることを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [13] The method for producing a high-strength hot-rolled steel sheet having excellent bendability, wherein the plating layer is a galvanization layer in [12].
［１４］前記［１２］において、前記めっき層が合金化亜鉛めっき層であることを特徴とする曲げ性に優れた高強度熱延鋼板の製造方法。 [14] The method for producing a high-strength hot-rolled steel sheet having excellent bendability, wherein the plating layer is an alloyed galvanized layer in [12].
　本発明によると、自動車の構造部材等の使途に好適な、引張強さ：980MPa以上であり且つ曲げ加工性に優れた高強度熱延鋼板が得られ、自動車部材の軽量化や自動車部材成形を可能とする等、その効果は著しく、また、高強度熱延鋼板の更なる用途展開が可能となり、産業上格段の効果を奏する。 According to the present invention, a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability suitable for the use of structural members of automobiles can be obtained. The effect is remarkable such as making it possible, and further application development of the high-strength hot-rolled steel sheet becomes possible, and there is a remarkable industrial effect.
　以下、本発明について詳細に説明する。
（高強度熱延鋼板）
　まず、本発明鋼板の組織及び炭化物の限定理由について説明する。本発明の熱延鋼板は、フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が8μm以下、該フェライト相の結晶粒内の炭化物の平均粒子径が10nm未満である組織を有する。 Hereinafter, the present invention will be described in detail.
(High-strength hot-rolled steel sheet)
First, the structure of the steel sheet of the present invention and the reasons for limiting the carbide will be described. The hot-rolled steel sheet of the present invention has a structure in which the area ratio of the ferrite phase is 95% or more, the average crystal grain size of the ferrite phase is 8 μm or less, and the average grain size of carbides in the ferrite phase crystal grains is less than 10 nm. Have.
フェライト相の面積率：95％以上
　熱延鋼板のマトリックスの金属組織は、加工性に優れたフェライト単相組織とすることが好ましい。ベイナイト相やマルテンサイト相、セメンタイト、パーライト等の第二相組織が鋼板組織に混入すると、硬度が互いに異なるフェライト相と第二相組織との界面でボイドが発生し、鋼板の曲げ加工性を低下させる。
　また、本発明では、鋼板中にTiやVなどの炭化物を析出させることで所望の鋼板強度を確保しようとするものであるが、これらの炭化物の多くは、熱延鋼板製造工程における仕上圧延終了後の冷却過程で、オーステナイトからフェライトへの変態と同時に相界面析出する炭化物である。したがって、炭化物をより多く得て所望の鋼板強度（引張強さ：980MPa以上）とするにはフェライト変態を促進させる必要があり、フェライト相の面積率が95％を下回ると引張強さを980MPa以上とすることが困難となる。 Area ratio of ferrite phase: 95% or more The matrix metal structure of the hot-rolled steel sheet is preferably a ferrite single-phase structure excellent in workability. When a second phase structure such as a bainite phase, martensite phase, cementite, or pearlite is mixed into the steel sheet structure, voids are generated at the interface between the ferrite phase and the second phase structure, which have different hardnesses, and the bending workability of the steel sheet is reduced. Let
In the present invention, it is intended to ensure the desired strength of the steel sheet by precipitating carbides such as Ti and V in the steel sheet, but most of these carbides are finished in the finish rolling in the hot-rolled steel sheet manufacturing process. It is a carbide that precipitates at the phase interface simultaneously with the transformation from austenite to ferrite in the subsequent cooling process. Therefore, it is necessary to promote ferrite transformation in order to obtain more carbide and achieve the desired steel sheet strength (tensile strength: 980 MPa or more). When the area ratio of the ferrite phase is less than 95%, the tensile strength is 980 MPa or more. It becomes difficult to do.
　以上の理由により、本発明では、熱延鋼板の金属組織をフェライト単相組織とすることが好ましいが、完全にフェライト単相でなくてもフェライト面積率が95％以上であれば、所望の強度（引張強さ：980MPa以上）が得られるため、フェライト相の面積率を95％以上とする。好ましくは98％以上である。 For the above reasons, in the present invention, it is preferable that the metal structure of the hot-rolled steel sheet is a ferrite single phase structure, but if the ferrite area ratio is 95% or more even if it is not completely a ferrite single phase, the desired strength (Tensile strength: 980 MPa or more) is obtained, so the area ratio of the ferrite phase is 95% or more. Preferably it is 98% or more.
　なお、本発明の熱延鋼板において、鋼板中に含有され得るフェライト相以外の組織としては、セメンタイト、パーライト、ベイナイト相、マルテンサイト相等が挙げられる。これらの組織が鋼板中に多量に存在すると、鋼板特性（曲げ加工性等）が低下する。そのため、これらの組織は極力低減することが好ましいが、鋼板の金属組織全体に対する合計面積率が5％以下であれば許容される。好ましくは2％以下である。 In the hot-rolled steel sheet of the present invention, examples of the structure other than the ferrite phase that can be contained in the steel sheet include cementite, pearlite, bainite phase, and martensite phase. When these structures are present in a large amount in the steel sheet, the steel sheet properties (bending workability, etc.) are deteriorated. Therefore, it is preferable to reduce these structures as much as possible, but it is acceptable if the total area ratio with respect to the entire metal structure of the steel sheet is 5% or less. Preferably it is 2% or less.
フェライト相の平均結晶粒径：8μm以下
　フェライト平均結晶粒径が8μmを上回ると、混粒組織（mixed grain size microstructure）となり易い。そして、混粒組織では、曲げ加工時に粗大なフェライト粒に応力が集中し易くなるため、鋼板の曲げ加工性が著しく低下する。したがって、フェライト相の平均結晶粒径の上限を8μmとする。好ましくは6μm以下であり、さらに好ましくは4.5μm以下である。 Average grain size of ferrite phase: 8 μm or less When the average grain size of ferrite exceeds 8 μm, it tends to be a mixed grain size microstructure. In the mixed grain structure, stress tends to concentrate on coarse ferrite grains during bending, so that the bending workability of the steel sheet is significantly reduced. Therefore, the upper limit of the average crystal grain size of the ferrite phase is 8 μm. Preferably it is 6 micrometers or less, More preferably, it is 4.5 micrometers or less.
フェライト結晶粒内の炭化物
　本発明の熱延鋼板は、強度を確保する点から、フェライト相の結晶粒内に炭化物を微細析出させることが必須となる。本発明においてフェライト相の結晶粒内に微細析出させる炭化物としては、Ti炭化物及びV炭化物及びTiとVの複合炭化物、或いはさらにNb、W、Mo、Hf、Zrを炭化物中に含むものが挙げられる。なお、これらの炭化物の多くは、熱延鋼板製造工程における仕上圧延終了後の冷却過程で、オーステナイトからフェライトへの変態と同時に相界面析出する炭化物である。 Carbides in ferrite crystal grains In the hot-rolled steel sheet of the present invention, it is essential to finely precipitate carbides in the ferrite phase crystal grains from the viewpoint of ensuring strength. Examples of the carbide finely precipitated in the ferrite phase grains in the present invention include Ti carbide, V carbide, and Ti and V composite carbide, or those containing Nb, W, Mo, Hf, and Zr in the carbide. . Most of these carbides are carbides that precipitate at the interface at the same time as the transformation from austenite to ferrite in the cooling process after finish rolling in the hot-rolled steel sheet manufacturing process.
フェライト結晶粒内の炭化物の平均粒子径：10nm未満
　本発明の熱延鋼板では、前記した炭化物、主にTi及びVの複合炭化物を微細に分散させることで強化を図っているが、炭化物が微細であるほど転位の運動を阻害する粒子数が増加するため、炭化物を分散することによって得られる強化量は増大する。したがって、本発明では、熱延鋼板を所望の引張強さ（980MPa）とする目的で、フェライト結晶粒内に分散させる炭化物の平均粒子径を10nm未満とする。好ましくは7nm未満であり、さらに好ましくは5nm以下である。 Average particle diameter of carbides in ferrite crystal grains: less than 10 nm In the hot rolled steel sheet of the present invention, the aforementioned carbides, mainly composite carbides of Ti and V, are strengthened by finely dispersing, but the carbides are fine. As the number of particles increases, the number of particles that inhibit dislocation movement increases, and the amount of strengthening obtained by dispersing carbide increases. Therefore, in the present invention, for the purpose of setting the hot-rolled steel sheet to a desired tensile strength (980 MPa), the average particle size of the carbide dispersed in the ferrite crystal grains is set to less than 10 nm. Preferably it is less than 7 nm, More preferably, it is 5 nm or less.
　次に、本発明熱延鋼板の成分組成の限定理由について説明する。なお、以下の成分組成を表す％は、特に断らない限り質量％（mass％）を意味するものとする。
C：0.06％以上0.1％以下
　Cは、TiやV、或いはさらにNbと結合し炭化物として鋼板中に微細分散（fine particle distribution）する。すなわちCは、微細な炭化物を形成してフェライト組織を著しく強化させる元素であり、熱延鋼板を強化するうえで必須の元素である。引張強さ980MPa以上の高強度鋼板を得るには、C含有量を少なくとも0.06％以上とする必要がある。一方、C含有量が0.1％を超えると、大量のセメンタイトが析出して鋼板の曲げ加工性が低下する。セメンタイトとマトリックス（フェライト）界面ではミクロボイドが生成し易く、このミクロボイドは鋼板の曲げ加工部での亀裂発生要因となるためである。したがって、C含有量を0.06％以上0.1％以下とする。好ましくは0.07％以上0.09％以下である。 Next, the reason for limiting the component composition of the hot-rolled steel sheet of the present invention will be described. In addition,% showing the following component composition shall mean the mass% (mass%) unless there is particular notice.
C: 0.06% or more and 0.1% or less C is combined with Ti, V, or Nb and is finely distributed in the steel sheet as carbides. That is, C is an element that forms fine carbides and remarkably strengthens the ferrite structure, and is an essential element for strengthening the hot-rolled steel sheet. In order to obtain a high-strength steel sheet having a tensile strength of 980 MPa or more, the C content needs to be at least 0.06%. On the other hand, when the C content exceeds 0.1%, a large amount of cementite precipitates and the bending workability of the steel sheet is lowered. This is because microvoids are easily generated at the cementite / matrix (ferrite) interface, and these microvoids cause cracks in the bent portion of the steel sheet. Therefore, the C content is set to 0.06% or more and 0.1% or less. Preferably they are 0.07% or more and 0.09% or less.
Si：0.09％以下
　Siは、延性（伸び）の低下をもたらすことなく鋼板強度を向上させる有効な元素として、従来の高強度鋼板では積極的に含有されている。しかしながら、Siは、鋼板表面に濃化し易く、鋼板表面にファイヤライト（Fe2SiO4）を形成する。このファイヤライトは鋼板表面に楔形となって形成するため、鋼板の曲げ加工の際、割れの起点となる。したがって、本発明ではSi含有量を極力低減することが望ましいが、0.09％までは許容できるため、Si含有量の上限を0.09％とする。好ましくは0.06％以下である。なお、Si含有量は不純物レベルまで低減してもよい。 Si: 0.09% or less Si is positively contained in conventional high-strength steel sheets as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, Si is easy to concentrate on the steel sheet surface, and forms firelite (Fe 2 SiO 4 ) on the steel sheet surface. Since this firelite is formed in a wedge shape on the surface of the steel sheet, it becomes a starting point of cracking when the steel sheet is bent. Therefore, in the present invention, it is desirable to reduce the Si content as much as possible, but 0.09% is acceptable, so the upper limit of the Si content is 0.09%. Preferably it is 0.06% or less. Note that the Si content may be reduced to the impurity level.
Mn：0.7％以上1.3％以下
　Mnは、フェライト相の結晶粒内に析出する炭化物を微細化するため、熱延鋼板の高強度化に有効な元素である。先述のとおり、本発明においてフェライト相の結晶粒内に析出する炭化物の多くは、熱延鋼板製造工程における仕上圧延終了後の冷却過程で、オーステナイトからフェライトへの変態と同時に相界面析出する。そのため、前記変態温度が高いと、炭化物が高温域で析出し、巻取りまでの冷却過程で炭化物が粗大化（coarsening）してしまう。 Mn: 0.7% or more and 1.3% or less Mn is an element effective for increasing the strength of a hot-rolled steel sheet in order to refine the carbides precipitated in the crystal grains of the ferrite phase. As described above, in the present invention, most of the carbides precipitated in the crystal grains of the ferrite phase are precipitated at the phase interface simultaneously with the transformation from austenite to ferrite in the cooling process after finishing rolling in the hot rolled steel sheet manufacturing process. For this reason, when the transformation temperature is high, carbides are precipitated in a high temperature region, and the carbides are coarsened in the cooling process until winding.
　このような問題に対し、Mnは鋼のオーステナイトからフェライトへの変態温度を低下させる作用を有するため、所定量のMnを含有させることで、前記変態温度を後述する巻取り温度域まで低下させ、鋼板の巻取りと同時に炭化物を析出させることができる。そして、このように高温域に長時間曝されることなく巻取りと同時に析出した炭化物は、微細な状態に維持される。炭化物を微細化して引張強さ980MPa以上の熱延鋼板を得るには、Mn含有量を少なくとも0.7％以上とする必要がある。一方、Mn含有量が1.3％を超えると、固溶Mnによる鋼板の加工性低下が顕著となるため、所望の曲げ加工性が得られない。したがって、Mn含有量は0.7％以上1.3％以下とする。好ましくは0.8％以上1.2％以下である。 For such problems, Mn has the effect of lowering the transformation temperature from austenite to ferrite of steel, so by containing a predetermined amount of Mn, the transformation temperature is lowered to the coiling temperature range described later, Carbide can be deposited simultaneously with the winding of the steel sheet. And the carbide | carbonized_material which precipitated simultaneously with winding, without being exposed to a high temperature range for a long time in this way is maintained in a fine state. In order to refine the carbide and obtain a hot rolled steel sheet having a tensile strength of 980 MPa or more, the Mn content needs to be at least 0.7%. On the other hand, if the Mn content exceeds 1.3%, the workability of the steel sheet due to solute Mn is significantly reduced, so that the desired bending workability cannot be obtained. Therefore, the Mn content is 0.7% or more and 1.3% or less. Preferably they are 0.8% or more and 1.2% or less.
P：0.03％以下
　Pは、粒界に偏析して加工時に粒界割れの起点となり、鋼板の曲げ加工性を劣化させる有害な元素であるため、極力低減することが好ましい。そこで、本発明では上記問題点を回避すべくP含有量を0.03％以下とする。好ましくは0.02％以下である。P含有量は不純物レベルまで低減してもよい。 P: 0.03% or less P is a harmful element that segregates at the grain boundary and becomes the starting point of grain boundary cracking during processing, and degrades the bending workability of the steel sheet, so it is preferably reduced as much as possible. Therefore, in the present invention, the P content is set to 0.03% or less in order to avoid the above problems. Preferably it is 0.02% or less. The P content may be reduced to the impurity level.
S：0.01％以下
　Sは、鋼中でMnSなどの介在物として存在する。この介在物は硬質であるため、鋼板の曲げ加工時にマトリックスと介在物との界面はボイドの起点となり、鋼板の曲げ加工性を低下させる。したがって、本発明では、S含有量を極力低減することが好ましく、0.01％以下とする。好ましくは0.008％以下である。Ｓ含有量はゼロであっても問題ない。 S: 0.01% or less S is present as an inclusion such as MnS in steel. Since the inclusions are hard, the interface between the matrix and the inclusions becomes a starting point of voids during bending of the steel sheet, thereby reducing the bending workability of the steel sheet. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, and set it to 0.01% or less. Preferably it is 0.008% or less. There is no problem even if the S content is zero.
Al：0.1％以下
　Alは、脱酸剤として作用する元素である。このような効果を得るためには0.02％以上含有することが望ましいが、Al含有量が0.1％を越えるとアルミナなどの介在物による曲げ加工性への悪影響が顕在化する。したがって、Al含有量は0.1％以下とする。好ましくは0.08％以下である。 Al: 0.1% or less Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is desirable to contain 0.02% or more. However, if the Al content exceeds 0.1%, an adverse effect on bending workability due to inclusions such as alumina becomes obvious. Therefore, the Al content is 0.1% or less. Preferably it is 0.08% or less.
N ：0.01％以下
　Nは、製鋼の段階で炭化物形成元素であるTiと結合して粗大なTi窒化物を形成し、微細な炭化物の形成を阻害するため著しく鋼板強度を低下させる。さらに、鋼板の曲げ加工時にマトリックスと粗大なTi窒化物の界面ではボイドが発生し易く、鋼板の曲げ加工性に悪影響をもたらす。したがって、N含有量は極力低減することが好ましく、0.01％以下とする。好ましくは0.008％以下である。Ｎ含有量はゼロであっても問題ない。 N: 0.01% or less N is combined with Ti, which is a carbide-forming element, at the steelmaking stage to form coarse Ti nitrides, which impede the formation of fine carbides and significantly reduce the steel sheet strength. Furthermore, voids are likely to occur at the interface between the matrix and coarse Ti nitride during bending of the steel sheet, which adversely affects the bending workability of the steel sheet. Therefore, the N content is preferably reduced as much as possible, and is 0.01% or less. Preferably it is 0.008% or less. There is no problem even if the N content is zero.
Ti：0.14％以上0.20％以下
　Tiは、Cと炭化物を形成して鋼板の高強度化に寄与する元素である。所望の熱延鋼板強度（引張強さ：980MPa以上）を確保するためには、Ti含有量を0.14％以上とする必要がある。一方、Ti含有量が0.20％を超えると、熱延鋼板を製造する際、熱間圧延前の鋼素材（スラブ）加熱によって粗大なTi炭化物を溶解することができず、最終的に得られる（巻取り後の）熱延鋼板に粗大なTi炭化物が残存する。このように粗大なTi炭化物が残存すると、熱延鋼板強度が低下するばかりか、マトリックスと粗大なTi炭化物との界面が鋼板曲げ加工時にボイドの起点となり、鋼板の曲げ加工性を低下させる。したがって、Ti含有量は0.14％以上0.20％以下とする。好ましくは0.15％以上0.19％以下である。 Ti: 0.14% or more and 0.20% or less Ti is an element that contributes to increasing the strength of steel sheets by forming carbides with C. In order to ensure the desired hot-rolled steel sheet strength (tensile strength: 980 MPa or more), the Ti content needs to be 0.14% or more. On the other hand, when the Ti content exceeds 0.20%, when manufacturing a hot-rolled steel sheet, coarse Ti carbide cannot be dissolved by heating the steel material (slab) before hot rolling, and finally obtained ( Coarse Ti carbide remains on the hot-rolled steel sheet after winding. If such coarse Ti carbide remains, not only the strength of the hot-rolled steel sheet is lowered, but also the interface between the matrix and the coarse Ti carbide serves as a starting point for voids during bending of the steel sheet, thereby reducing the bending workability of the steel sheet. Therefore, the Ti content is 0.14% or more and 0.20% or less. Preferably it is 0.15% or more and 0.19% or less.
V ：0.07％以上0.14％以下
　Vは、Tiと同様、Cと炭化物を形成して鋼板の高強度化に寄与する元素である。Vは、Tiと結合して微細な複合炭化物を形成するため、鋼板の高強度化に有効である。所望の熱延鋼板強度（引張強さ：980MPa以上）を確保するためには、V含有量を0.07％以上とする必要がある。一方、Ti含有量よりもV含有量が多くなると、Vは析出し難い状態となり固溶状態として鋼板中に残存するV量が多くなる。固溶状態にあるVは鋼板の曲げ加工性を劣化させるため、V含有量はTi含有量以下、すなわち0.14％以下とする必要がある。したがって、V含有量は0.07％以上0.14％以下とする。好ましくは0.08％以上0.13％以下である。 V: 0.07% or more and 0.14% or less V, like Ti, is an element that contributes to increasing the strength of steel sheets by forming carbides with C. V is effective for increasing the strength of the steel sheet because it combines with Ti to form a fine composite carbide. In order to ensure the desired hot-rolled steel sheet strength (tensile strength: 980 MPa or more), the V content needs to be 0.07% or more. On the other hand, when the V content is higher than the Ti content, V becomes difficult to precipitate and the amount of V remaining in the steel sheet as a solid solution state increases. Since V in a solid solution state deteriorates the bending workability of the steel sheet, the V content needs to be Ti content or less, that is, 0.14% or less. Therefore, the V content is 0.07% or more and 0.14% or less. Preferably they are 0.08% or more and 0.13% or less.
Nb：0.01％以上0.05％以下
　以上が、本発明における基本成分であるが、上記した基本成分に加えてさらにNb:0.01％以上0.05％以下を含有してもよい。
　Nbは、実質的にフェライト単相組織を有する熱延鋼板を製造する際の熱間圧延工程において、オーステナイトからフェライトへの変態前のオーステナイト粒の再結晶を阻害して未再結晶組織とする作用がある。未再結晶組織では、熱間圧延によるひずみエネルギーが蓄積し易く、フェライト相の核生成サイト数が増大する。そのため、Nbを添加すると、熱間圧延工程においてフェライト相の核生成サイト数が増大する結果、フェライト相の結晶粒を微細化することができる。このような効果を得るには、Nb含有量を0.01％以上とすることが好ましい。一方、熱間圧延工程で鋼に過度に歪みエネルギーを加えると、オーステナイト→フェライト変態温度が上昇するため、微細な炭化物が得られなくなるおそれがある。このような観点から、Nb含有量は0.05％以下とすることが好ましい。より好ましくは0.02％以上0.04％以下である。 Nb: 0.01% or more and 0.05% or less The above is the basic component in the present invention. In addition to the above basic components, Nb: 0.01% or more and 0.05% or less may be further contained.
Nb has the effect of inhibiting the recrystallization of austenite grains before transformation from austenite to ferrite in the hot rolling process when producing a hot rolled steel sheet having a ferrite single-phase structure substantially to make an unrecrystallized structure There is. In the non-recrystallized structure, strain energy due to hot rolling is easily accumulated, and the number of nucleation sites of the ferrite phase increases. Therefore, when Nb is added, the number of nucleation sites of the ferrite phase is increased in the hot rolling process, so that the crystal grains of the ferrite phase can be refined. In order to obtain such an effect, the Nb content is preferably set to 0.01% or more. On the other hand, if strain energy is excessively applied to the steel in the hot rolling process, the austenite → ferrite transformation temperature rises, and fine carbides may not be obtained. From such a viewpoint, the Nb content is preferably 0.05% or less. More preferably, it is 0.02% or more and 0.04% or less.
Mo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下
　また、前記基本成分に加えてさらにMo、W、Zr及びHfのうちのいずれか１種以上を含有する場合、これらの含有量はMo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下に制限することが好ましい。
　Mo、W、Zr及びHfは、炭化物を形成して鋼板の高強度化に寄与する元素であるが、固溶状態として残存する量も多い。これらの固溶元素はマトリックスの加工性を低下させ、鋼板の曲げ加工性に悪影響を及ぼす。Mo、W、Zr、Hfは、含有量に対して析出する割合が低く、固溶元素として鋼板中に残存する量が多い。そのため、これらの含有量は極力低減することが望ましいが、それぞれ0.05％まで許容できるため、上限量を0.05％とした。好ましくは0.03％以下である。なお、Mo、W、Zr、Hfの含有量はゼロとしてもよい。 Mo: 0.05% or less, W: 0.05% or less, Zr: 0.05% or less, Hf: 0.05% or less In addition to the basic components, one or more of Mo, W, Zr and Hf are further contained. In such a case, it is preferable to limit these contents to Mo: 0.05% or less, W: 0.05% or less, Zr: 0.05% or less, and Hf: 0.05% or less.
Mo, W, Zr, and Hf are elements that form carbides and contribute to increasing the strength of the steel sheet, but a large amount remains as a solid solution. These solid solution elements deteriorate the workability of the matrix and adversely affect the bending workability of the steel sheet. Mo, W, Zr, and Hf have a low rate of precipitation with respect to the content, and a large amount remains in the steel sheet as a solid solution element. For this reason, it is desirable to reduce these contents as much as possible, but 0.05% is acceptable, so the upper limit was set to 0.05%. Preferably it is 0.03% or less. Note that the contents of Mo, W, Zr, and Hf may be zero.
その他の含有成分
　また、上述した基本成分に加えてさらに、O（酸素）、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、B、Ni、Cr、Sb、Cu、Sn、Mg及びCaのうちのいずれか１種以上を合計で0.2％以下含有してもよい。これらの元素は、鋼板の曲げ加工性の観点から合計で0.2％までは許容できる。好ましくは合計で0.09％以下とする。上記以外の成分は、Fe及び不可避的不純物である。 Other components In addition to the basic components described above, O (oxygen), Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, One or more of Pd, Pt, Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, B, Ni, Cr, Sb, Cu, Sn, Mg, and Ca You may contain 0.2% or less in total. These elements are acceptable up to 0.2% in total from the viewpoint of bending workability of the steel sheet. Preferably, the total content is 0.09% or less. Components other than the above are Fe and inevitable impurities.
　また、本発明の熱延鋼板の表面にめっき層を形成してもよい。表面にめっき層を形成することにより、熱延鋼板の耐食性が向上し、厳しい腐食環境下で使用される自動車部品などへの適用が可能になる。
　めっき層の種類は特に問わず、電気めっき層、無電解めっき層のいずれも適用可能である。また、めっき層の合金成分も特に問わず、溶融亜鉛めっき層、合金化溶融亜鉛めっき層などが好適な例として挙げられるが、勿論、これらに限定されず従前公知のものがいずれも適用可能である。 Moreover, you may form a plating layer in the surface of the hot-rolled steel plate of this invention. By forming a plating layer on the surface, the corrosion resistance of the hot-rolled steel sheet is improved, and it becomes possible to apply it to automobile parts used in severe corrosive environments.
The type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied. Further, the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer and the like can be mentioned as suitable examples, but of course, it is not limited to these, and any conventionally known one can be applied. is there.
（高強度熱延鋼板の製造方法）
　次に、本発明の熱延鋼板の製造方法について説明する。
本発明は、上記した組成の鋼素材（鋼スラブ）を加熱し、粗圧延と仕上圧延からなる熱間圧延を施し、仕上圧延終了後、冷却し、巻き取り、熱延鋼板とする。
　そして、前記鋼素材の加熱温度を1100℃以上1350℃以下とし、前記仕上圧延の仕上圧延温度を850℃以上とし、前記冷却を仕上圧延終了後から3秒以内に開始し、前記冷却の平均冷却速度を20℃/s以上とし、前記巻き取りの巻取り温度を550℃以上700℃以下とすることを特徴とする。 (Method for producing high-strength hot-rolled steel sheet)
Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
In the present invention, a steel material (steel slab) having the above composition is heated, subjected to hot rolling consisting of rough rolling and finish rolling, cooled after completion of finish rolling, and wound into a hot rolled steel sheet.
Then, the heating temperature of the steel material is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, and the cooling is started within 3 seconds after finishing the finish rolling, and the average cooling of the cooling The speed is 20 ° C./s or more, and the winding temperature of the winding is 550 ° C. or more and 700 ° C. or less.
　本発明において、鋼の溶製方法は特に限定されず、転炉、電気炉等、公知の溶製方法を採用することができる。また、真空脱ガス炉にて2次精錬を行ってもよい。その後、生産性や品質上の問題から連続鋳造法によりスラブ（鋼素材）とするのが好ましいが、造塊－分塊圧延法、薄スラブ連鋳法等、公知の鋳造方法でスラブとしても良い。 In the present invention, 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. .
鋼素材の加熱温度：1100℃以上1350℃以下
　上記の如く得られた鋼素材（鋼スラブ）に、粗圧延及び仕上圧延を施すが、本発明においては、粗圧延に先立ち鋼素材を加熱して実質的に均質なオーステナイト相とし、粗大な炭化物を溶解する必要がある。鋼素材の加熱温度が1100℃を下回ると、粗大な炭化物が溶解しないため、熱間圧延終了後の冷却・巻取り工程で微細分散する炭化物の量が減じることとなり、最終的に得られる熱延鋼板の強度が著しく低下する。一方、上記加熱温度が1350℃を上回ると、スケールが噛み込み、鋼板の表面性状（surface appearance quality）を悪化させる。 Heating temperature of steel material: 1100 ° C or higher and 1350 ° C or lower The steel material (steel slab) obtained as described above is subjected to rough rolling and finish rolling. In the present invention, the steel material is heated prior to rough rolling. It is necessary to form a substantially homogeneous austenite phase and dissolve coarse carbides. When the heating temperature of the steel material is below 1100 ° C, coarse carbides do not dissolve, so the amount of carbides that are finely dispersed in the cooling and winding process after hot rolling is reduced, and the hot rolling finally obtained is reduced. The strength of the steel sheet is significantly reduced. On the other hand, when the heating temperature exceeds 1350 ° C., the scale bites and deteriorates the surface appearance quality of the steel sheet.
　以上の理由により、鋼素材の加熱温度は1100℃以上1350℃以下とする。好ましくは1150℃以上1320℃以下である。但し、鋼素材に熱間圧延を施すに際し、鋳造後の鋼素材が1100℃以上1350℃以下の温度域にある場合、或いは鋼素材の炭化物が溶解している場合には、鋼素材を加熱することなく直送圧延してもよい。なお、粗圧延条件については特に限定されない。 For the above reasons, the heating temperature of the steel material is set to 1100 ° C to 1350 ° C. Preferably they are 1150 degreeC or more and 1320 degrees C or less. However, when hot rolling the steel material, if the steel material after casting is in the temperature range of 1100 ° C or higher and 1350 ° C or lower, or if the carbide of the steel material is dissolved, the steel material is heated. Direct rolling may be performed without any problem. The rough rolling conditions are not particularly limited.
仕上圧延温度：850℃以上
　仕上圧延温度が850℃を下回ると、仕上圧延中にフェライト変態が開始してフェライト粒が伸展された組織となるうえ、部分的にフェライト粒が成長した混粒組織となるため、熱延鋼板の曲げ加工性が著しく低下する。したがって、仕上圧延温度は850℃以上とする。好ましくは870℃以上である。なお、仕上圧延温度の上限は特に定めないが、仕上圧延温度はスラブ再加熱温度や通板速度、鋼板板厚により決定される。そのため、実質的に仕上圧延温度の上限は990℃以下である。 Finish rolling temperature: 850 ° C or more When the finish rolling temperature is lower than 850 ° C, ferrite transformation starts during finish rolling, and the ferrite grains are expanded, and a mixed grain structure in which ferrite grains partially grow Therefore, the bending workability of the hot-rolled steel sheet is significantly reduced. Therefore, the finish rolling temperature is 850 ° C. or higher. Preferably it is 870 degreeC or more. The upper limit of the finish rolling temperature is not particularly defined, but the finish rolling temperature is determined by the slab reheating temperature, the sheet feed speed, and the steel plate thickness. Therefore, the upper limit of the finish rolling temperature is substantially 990 ° C. or less.
仕上圧延終了後から強制冷却を開始するまでの時間：3秒以内
　仕上圧延直後の高温状態の鋼板においては、オーステナイト相に蓄積されたひずみエネルギーが大きいため、ひずみ誘起析出による炭化物が生じる。この炭化物は、高温で析出するため粗大化し易いことから、ひずみ誘起析出が生じると微細な析出物が得られ難くなる。したがって、本発明では、ひずみ誘起析出を抑制する目的で熱間圧延終了後速やかに強制冷却を開始する必要があり、仕上圧延終了後から、遅くとも3秒以内に冷却を開始する。好ましくは2秒以内である。 Time from the end of finish rolling to the start of forced cooling: within 3 seconds In a high-temperature steel sheet immediately after finish rolling, carbides are generated by strain-induced precipitation because the strain energy accumulated in the austenite phase is large. Since this carbide precipitates at a high temperature and is easy to coarsen, it is difficult to obtain a fine precipitate when strain-induced precipitation occurs. Therefore, in the present invention, for the purpose of suppressing strain-induced precipitation, it is necessary to start forced cooling immediately after the end of hot rolling, and cooling is started within 3 seconds at the latest after finishing rolling. Preferably it is within 2 seconds.
平均冷却速度：20℃/s以上
　上記のとおり、仕上圧延終了後の鋼板の高温に維持される時間が長いほど、ひずみ誘起析出による炭化物の粗大化が進行し易くなる。また、本発明においては、鋼板に所定量のMnを含有させることよってオーステナイト→フェライト変態を抑制しているものの、冷却速度が小さいと高温でフェライト変態が開始し、炭化物が粗大化し易くなる。そのため、仕上圧延後は急冷する必要があり、上記問題を回避するには20℃/s以上の平均冷却速度で冷却する必要がある。好ましい平均冷却速度は、40℃/s以上である。ただし、仕上圧延終了後の冷却速度が過剰に大きくなると、巻取温度の制御が困難となり安定した熱延鋼板強度が得られ難くなることが懸念されるため、150℃/s以下とすることが好ましい。 Average cooling rate: 20 ° C./s or more As described above, the longer the time that the steel sheet after finish rolling is maintained at a high temperature, the more easily the coarsening of the carbide by strain-induced precipitation proceeds. Further, in the present invention, the austenite → ferrite transformation is suppressed by containing a predetermined amount of Mn in the steel sheet, but if the cooling rate is low, the ferrite transformation starts at a high temperature, and the carbide tends to become coarse. Therefore, it is necessary to rapidly cool after finish rolling, and in order to avoid the above problem, it is necessary to cool at an average cooling rate of 20 ° C./s or more. A preferable average cooling rate is 40 ° C./s or more. However, if the cooling rate after finishing rolling is excessively increased, it is difficult to control the coiling temperature and it is difficult to obtain stable hot-rolled steel sheet strength. preferable.
巻取り温度：550℃以上700℃以下
　巻取り温度が550℃を下回ると十分な量の炭化物が得られず、鋼板強度が低下する。一方、巻取り温度が700℃を超えると、析出した炭化物が粗大化するため鋼板強度が低下する。したがって、巻取り温度の範囲は550℃以上700℃以下とする。好ましくは580℃以上680℃以下である。 Winding temperature: 550 ° C. or higher and 700 ° C. or lower If the winding temperature is lower than 550 ° C., a sufficient amount of carbide cannot be obtained, and the steel sheet strength decreases. On the other hand, when the coiling temperature exceeds 700 ° C., the precipitated carbide is coarsened, so that the steel sheet strength is lowered. Accordingly, the coiling temperature range is 550 ° C. or higher and 700 ° C. or lower. Preferably they are 580 degreeC or more and 680 degrees C or less.
　なお、熱間圧延した巻き取り後の熱延鋼板は、表面にスケールが付着した状態であっても、酸洗を行うことによりスケールを除去した状態であっても、その特性が変わることはなく、いずれの状態においても前記した優れた特性を発現する。また、本発明では、巻き取り後の熱延鋼板にめっき処理を施して、熱延鋼板表面にめっき層を形成してもよい。 Note that the hot-rolled steel sheet after being rolled by hot rolling does not change its properties even if the scale is attached to the surface or the scale is removed by pickling. The above-described excellent characteristics are exhibited in any state. In the present invention, the hot-rolled steel sheet after winding may be plated to form a plating layer on the surface of the hot-rolled steel sheet.
　本発明の熱延鋼板は、740℃までの加熱処理を短時間施しても材質変動が小さい。そのため、鋼板に耐食性を付与する目的で、本発明の熱延鋼板にめっき処理を施し、その表面にめっき層を具えることができる。めっき処理における加熱温度は740℃以下でも製造可能であることから、本発明の熱延鋼板にめっき処理を施しても前記した本発明の効果を損なうことはない。めっき層の種類は特に問わず、電気めっき層、無電解めっき層のいずれも適用可能である。また、めっき層の合金成分も特に問わず、溶融亜鉛めっき層、合金化溶融亜鉛めっき層などが好適な例として挙げられるが、勿論、これらに限定されず従前公知のものがいずれも適用可能である。 The hot-rolled steel sheet of the present invention has little material fluctuation even when subjected to heat treatment up to 740 ° C. for a short time. Therefore, for the purpose of imparting corrosion resistance to the steel sheet, the hot-rolled steel sheet of the present invention can be plated, and a plating layer can be provided on the surface thereof. Since it can be manufactured even at a heating temperature of 740 ° C. or lower in the plating process, the effect of the present invention described above is not impaired even if the hot-rolled steel sheet of the present invention is plated. The type of the plating layer is not particularly limited, and any of an electroplating layer and an electroless plating layer can be applied. Further, the alloy component of the plating layer is not particularly limited, and a hot dip galvanized layer, an alloyed hot dip galvanized layer and the like can be mentioned as suitable examples, but of course, it is not limited to these, and any conventionally known one can be applied. is there.
　めっき処理の方法も特に問わず、従前公知の方法がいずれも適用可能である。例えば、焼鈍温度を740℃以下とした連続めっきライン（continuous galvanizing/galvannealing line）に通板させたのち、めっき浴に鋼板を浸漬して引き上げる方法などが挙げられる。また、めっき処理後にガス炉などの炉内で鋼板表面を加熱して合金化処理を施してもよい。 The plating method is not particularly limited, and any conventionally known method can be applied. For example, after passing through a continuous plating line (continuous galvanizing / galvannealing line) with an annealing temperature of 740 ° C. or lower, a method of immersing and pulling up the steel plate in a plating bath may be used. Further, after the plating treatment, the steel plate surface may be heated in a furnace such as a gas furnace to perform the alloying treatment.
　以上のように、本発明によると、鋼板組成及び製造条件を適正化することで、フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が8μm以下であり、前記フェライト相の結晶粒内の炭化物平均粒子径が10nm未満である組織を有する熱延鋼板が得られる。また、本発明では、曲げ加工性を高める目的で鋼板中に存在する固溶元素及び粗大な介在物を低減したうえで高強度化を図っているため、曲げ加工性に優れた高強度熱延鋼板とすることができる。 As described above, according to the present invention, by optimizing the steel sheet composition and production conditions, the ferrite phase area ratio is 95% or more, the average grain size of the ferrite phase is 8 μm or less, the ferrite phase of the ferrite phase A hot-rolled steel sheet having a structure in which the carbide average particle diameter in the crystal grains is less than 10 nm is obtained. Further, in the present invention, in order to increase the bending workability, the solid solution elements and coarse inclusions present in the steel sheet are reduced, and the strength is increased, so that the high strength hot rolling excellent in bending workability is achieved. It can be a steel plate.
　さらに、本発明では、鋼板に含まれる炭化物形成元素（Ti及びV、或いはさらにNb、W、Mo、Hf、Zr）の含有量を最適化したうえ、熱延鋼板の製造条件を規定している。これにより、フェライト結晶粒内に上記した平均粒子径が10nm未満である炭化物を十分に析出させることができ、優れた曲げ加工性を維持しつつ熱延鋼板の引張強さを980MPa以上にまで高めることができる。なお、本発明は引張強さ1100MPa以下の熱延鋼板に適用することが好ましく、引張強さ1052 MPa以下の熱延鋼板に適用することがより好ましい。 Furthermore, in the present invention, the content of carbide forming elements (Ti and V, or even Nb, W, Mo, Hf, Zr) contained in the steel sheet is optimized, and the production conditions for the hot rolled steel sheet are specified. . As a result, the above-mentioned carbide having an average particle diameter of less than 10 nm can be sufficiently precipitated in the ferrite crystal grains, and the tensile strength of the hot-rolled steel sheet is increased to 980 MPa or more while maintaining excellent bending workability. be able to. The present invention is preferably applied to a hot rolled steel sheet having a tensile strength of 1100 MPa or less, and more preferably applied to a hot rolled steel sheet having a tensile strength of 1052 MPa or less.
　表１に示す組成を有する肉厚250mmの鋼素材（鋼スラブ）に、表２に示す熱延条件で熱間圧延を施して板厚1.4～3.2mmの熱延鋼板とした。なお、表２に記載の冷却速度は、仕上圧延温度から巻取り温度までの平均冷却速度である。
　また、得られた熱延鋼板の一部に対して、焼鈍温度720℃の溶融亜鉛めっきラインに通板し、その後、460℃のめっき浴（めっき組成：Zn－0.13mass％Al）に浸漬し、溶融亜鉛めっき材（GI材）とした。また一部溶融亜鉛めっき材（GI材）に対しては、溶融亜鉛めっきラインに通板、めっき浴浸漬に次いで、520℃で合金化処理を施して合金化溶融亜鉛めっき材（GA材）とした。めっき付着量はGI材、GA材ともに片面当たり45～55g/m2とした。
　なお、鋼板No.3～5、12～18を除き、巻き取りまでの冷却中にオーステナイトからフェライトへの変態は生じていないことを、別途確認している。 A steel material (steel slab) having a composition shown in Table 1 and having a thickness of 250 mm was hot rolled under the hot rolling conditions shown in Table 2 to obtain a hot rolled steel sheet having a thickness of 1.4 to 3.2 mm. In addition, the cooling rate described in Table 2 is an average cooling rate from the finish rolling temperature to the winding temperature.
In addition, a part of the obtained hot-rolled steel sheet is passed through a hot dip galvanizing line with an annealing temperature of 720 ° C, and then immersed in a 460 ° C plating bath (plating composition: Zn-0.13 mass% Al). A hot-dip galvanized material (GI material) was used. In addition, for partially hot dip galvanized materials (GI materials), passing through the hot dip galvanizing line and immersion in the plating bath, followed by alloying treatment at 520 ° C, and galvannealed galvanized material (GA material) did. The amount of plating applied was 45 to 55 g / m 2 per side for both GI and GA materials.
Except for steel plates Nos. 3 to 5 and 12 to 18, it was confirmed separately that transformation from austenite to ferrite did not occur during cooling until winding.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
　上記により得られた熱延鋼板（熱延鋼板、GI材、GA材）から試験片を採取し、組織観察、引張試験、曲げ試験を行い、フェライト相の面積率、フェライト相以外の組織の種類及び面積率、フェライト相の平均結晶粒径、炭化物の平均粒子径、降伏強度、引張強さ、伸び、限界曲げ半径を求めた。試験方法は次のとおりとした。 Samples are taken from the hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material) obtained as described above, and subjected to structure observation, tensile test, and bending test. The area ratio of the ferrite phase and the type of structure other than the ferrite phase The area ratio, the average crystal grain size of the ferrite phase, the average grain size of the carbide, the yield strength, the tensile strength, the elongation, and the critical bending radius were obtained. The test method was as follows.
（ｉ）組織観察
　フェライト相の面積率は以下の手法により評価した。圧延方向に平行な断面の板厚中心部を、5％ナイタール（nital）による腐食現出組織を走査型光学顕微鏡で400倍に拡大して10視野分撮影した。フェライト相は、粒内に腐食痕やセメンタイトが観察されない形態を有する組織である。また、ポリゴナルフェライト、ベイニティックフェライト、アシキュラーフェライト及びグラニュラーフェライトをフェライトとして、フェライト相の面積率、フェライト相の平均粒径及びフェライト相の結晶粒内の炭化物の平均粒径を導出した。
　フェライト相の面積率は、画像解析によりフェライト相とベイナイトやマルテンサイト等のフェライト相以外を分離し、観察視野に対するフェライト相の面積率によって求めた。このとき、線状の形態として観察される粒界はフェライト相の一部として計上した。得られたフェライト相の面積率を表３に示す。
　フェライト相の平均結晶粒径は、上記400倍に拡大して撮影し代表的な写真3枚について水平線及び垂直線をそれぞれ3本ずつ引きASTM E 112-10に準拠した切断法によって求め、3枚の平均値を最終的な平均結晶粒径とした。得られた平均結晶粒径を表３に示す。
　フェライト相の結晶粒内の炭化物の平均粒子径は、得られた熱延鋼板の板厚中央部から薄膜法によってサンプルを作製し、透過型電子顕微鏡（倍率：135000倍）で観察を行い、100点以上の析出物粒子径の平均によって求めた。この析出物粒子径を算出する上で、粒子径が1.0μm以上の粗大なセメンタイトや窒化物は含まないものとした。得られた炭化物の平均粒径を表３に示す。 (I) Structure observation The area ratio of the ferrite phase was evaluated by the following method. The central portion of the plate thickness in the cross section parallel to the rolling direction was photographed for 10 fields of view by corroding the appearance of corrosion by 5% nital 400 times with a scanning optical microscope. The ferrite phase is a structure having a form in which corrosion marks and cementite are not observed in the grains. Further, using polygonal ferrite, bainitic ferrite, acicular ferrite, and granular ferrite as ferrite, the area ratio of the ferrite phase, the average grain size of the ferrite phase, and the average grain size of the carbide in the crystal grains of the ferrite phase were derived.
The area ratio of the ferrite phase was obtained by separating the ferrite phase from the ferrite phase other than the ferrite phase such as bainite and martensite by image analysis, and obtaining the area ratio of the ferrite phase with respect to the observation field. At this time, the grain boundary observed as a linear form was counted as a part of the ferrite phase. Table 3 shows the area ratio of the obtained ferrite phase.
The average crystal grain size of the ferrite phase was obtained by enlarging the above 400 times and taking three representative photographs by drawing three horizontal lines and three vertical lines, respectively, by a cutting method in accordance with ASTM E 112-10. Was the final average crystal grain size. Table 3 shows the average crystal grain size obtained.
The average particle size of the carbides in the ferrite phase grains was measured using a transmission electron microscope (magnification: 135000 times) by preparing a sample from the center of the thickness of the obtained hot rolled steel sheet using a thin film method. It calculated | required by the average of the particle diameter of the precipitate more than a point. In calculating the particle size of the precipitate, coarse cementite and nitride having a particle size of 1.0 μm or more were not included. Table 3 shows the average particle size of the obtained carbide.
（ｉｉ）引張試験
　得られた熱延鋼板から圧延方向と垂直方向にJIS5号引張試験片を作製し、JIS Z 2241(2011)の規定に準拠した引張試験を5回行い、平均の降伏強度(YS)、引張強さ(TS)、全伸び(El)を求めた。なお、引張試験のクロスヘッドスピードは10mm/minとした。 (Ii) Tensile test A JIS No. 5 tensile test piece was produced from the obtained hot-rolled steel sheet in the direction perpendicular to the rolling direction, the tensile test was conducted 5 times in accordance with the provisions of JIS Z 2241 (2011), and the average yield strength ( YS), tensile strength (TS), and total elongation (El) were determined. The crosshead speed in the tensile test was 10 mm / min.
（ｉｉｉ）曲げ試験（曲げ加工性評価）
　得られた熱延鋼板から、長手方向が圧延方向に対して直角になるように短冊状の試験片（100W×35L mm）をせん断加工によって採取した。このとき、せん断面と破断面が試験片端面で同一の方向とした。以上のようにして採取した試験片を用いてJIS Z 2248に準拠したVブロック法による曲げ試験を3回行い、試験後サンプルの湾曲部外側を肉眼で観察し、裂けや疵等の欠点がないものを合格した。種々の内径半径を有する押金具を用いて試験を行い、下式に示すように、合格となった押金具の最小内側半径R（mm）を熱延鋼板の板厚t（mm）で除した値（R/t）を限界曲げ半径とした。
（限界曲げ半径）＝（合格となった押金具の最小内側半径R）/（鋼板板厚t）
限界曲げ半径は小さい値であるほど良い結果であることを示す。限界曲げ半径が2.0以下である場合の評価を良好“○”とし、限界曲げ半径が2.0超である場合の評価を不良“×”とした。
　以上により得られた結果を表３に示す。 (Iii) Bending test (bending workability evaluation)
From the obtained hot-rolled steel sheet, a strip-shaped test piece (100 W × 35 L mm) was collected by shearing so that the longitudinal direction was perpendicular to the rolling direction. At this time, the shear plane and the fracture surface were in the same direction on the end face of the test piece. Using the test piece collected as described above, the bending test by the V-block method in accordance with JIS Z 2248 is performed three times, and the outside of the curved part of the sample after the test is observed with the naked eye, and there are no defects such as tears and wrinkles. Passed things. Tests were conducted using metal fittings with various inner radiuses, and the minimum inner radius R (mm) of the accepted metal fittings was divided by the thickness t (mm) of the hot-rolled steel sheet, as shown in the following formula. The value (R / t) was taken as the critical bending radius.
(Limit bending radius) = (Minimum inner radius R of the accepted metal fittings) / (Steel plate thickness t)
A smaller limit bending radius indicates a better result. The evaluation when the critical bending radius was 2.0 or less was evaluated as “good”, and the evaluation when the critical bending radius was over 2.0 was evaluated as “poor”.
The results obtained as described above are shown in Table 3.
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
　本発明例はいずれも、引張強さTS：980MPa以上であり且つ曲げ加工性にも優れ、強度と加工性を兼備した熱延鋼板となっている。一方、本発明の範囲を外れる比較例は、所定の高強度が確保できていないか、良好な曲げ加工性が得られていないことがわかった。 All examples of the present invention are hot-rolled steel sheets having a tensile strength of TS: 980 MPa or more, excellent bending workability, and both strength and workability. On the other hand, it was found that the comparative example out of the scope of the present invention did not ensure a predetermined high strength or did not obtain good bending workability.
　本発明によれば、自動車の構造部材等の使途に好適な、引張強さ：980MPa以上であり且つ曲げ加工性に優れた高強度熱延鋼板が得られ、自動車部材の軽量化と自動車部材成形との両立が可能となる。
  According to the present invention, a high-strength hot-rolled steel sheet having a tensile strength of 980 MPa or more and excellent bending workability suitable for the use of structural members of automobiles can be obtained. It becomes possible to achieve both.

Claims (14)

  1. 　質量％で、
    C：0.06％以上0.1％以下、
    Si：0.09％以下、
    Mn：0.7％以上1.3％以下、
    P：0.03％以下、
    S：0.01％以下、
    Al：0.1％以下、
    N：0.01％以下、
    Ti：0.14％以上0.20％以下
    V：0.07％以上0.14％以下
    を含有し、残部がFe及び不可避的不純物である組成からなり、
    　フェライト相の面積率が95％以上、該フェライト相の平均結晶粒径が8μm以下、該フェライト相の結晶粒内の炭化物の平均粒子径が10nm未満である組織を有し、引張強さが980MPa以上である高強度熱延鋼板。     % By mass
    C: 0.06% to 0.1%,
    Si: 0.09% or less,
    Mn: 0.7% or more and 1.3% or less,
    P: 0.03% or less,
    S: 0.01% or less,
    Al: 0.1% or less,
    N: 0.01% or less,
    Ti: 0.14% to 0.20%
    V: Containing 0.07% or more and 0.14% or less, with the balance being Fe and inevitable impurities,
    The ferrite phase has an area ratio of 95% or more, the ferrite phase has an average crystal grain size of 8 μm or less, and the ferrite grain has an average grain size of less than 10 nm, and a tensile strength of 980 MPa. This is the high-strength hot-rolled steel sheet.
  2. 　前記組成に加えてさらに、質量％でNb：0.01％以上0.05％以下を含有する請求項１に記載の高強度熱延鋼板。 The high-strength hot-rolled steel sheet according to claim 1, further comprising Nb: 0.01% or more and 0.05% or less by mass% in addition to the composition.
  3. 　前記組成に加えてさらに、Mo、W、Zr、Hfのいずれか１種以上を含有し、これらの含有量を、Mo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下に制限してなる請求項１又は２に記載の高強度熱延鋼板。 In addition to the above composition, the composition further contains at least one of Mo, W, Zr, and Hf. The contents of Mo, 0.05% or less, W: 0.05% or less, Zr: 0.05% or less, Hf: The high-strength hot-rolled steel sheet according to claim 1 or 2, which is limited to 0.05% or less.
  4. 　前記組成に加えてさらに、質量％で、O（酸素）、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、B、Ni、Cr、Sb、Cu、Sn、Mg及びCaのうちの１種以上を合計で0.2％以下含有する請求項１～３のいずれかに記載の高強度熱延鋼板。 In addition to the above composition, in mass%, O (oxygen), Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, Pd, Pt, Contains a total of 0.2% or less of one or more of Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, B, Ni, Cr, Sb, Cu, Sn, Mg and Ca The high-strength hot-rolled steel sheet according to any one of claims 1 to 3.
  5. 　前記鋼板表面に、めっき層をさらに有する請求項１～４のいずれかに記載の高強度熱延鋼板。 The high-strength hot-rolled steel sheet according to any one of claims 1 to 4, further comprising a plating layer on the steel sheet surface.
  6. 　前記めっき層が亜鉛めっき層である請求項５に記載の高強度熱延鋼板。 The high-strength hot-rolled steel sheet according to claim 5, wherein the plated layer is a galvanized layer.
  7. 　前記めっき層が合金化亜鉛めっき層である請求項５に記載の高強度熱延鋼板。 The high-strength hot-rolled steel sheet according to claim 5, wherein the plated layer is an alloyed galvanized layer.
  8. 　鋼素材を加熱し、粗圧延と仕上圧延からなる熱間圧延を施し、仕上圧延終了後、冷却し、巻き取り、熱延鋼板とするにあたり、
    　前記鋼素材を、質量％で、
    C：0.06％以上0.1％以下、
    Si：0.09％以下、
    Mn：0.7％以上1.3％以下、
    P：0.03％以下、
    S：0.01％以下、
    Al：0.1％以下、
    N：0.01％以下、
    Ti：0.14％以上0.20％以下
    V：0.07％以上0.14％以下
    を含有し、残部がFe及び不可避的不純物からなる組成とし、
    　前記鋼素材の加熱温度を1100℃以上1350℃以下とし、前記仕上圧延の仕上圧延温度を850℃以上とし、前記冷却を仕上圧延終了後から3秒以内に開始し、前記冷却の平均冷却速度を20℃/s以上とし、前記巻き取りの巻取り温度を550℃以上700℃以下とする高強度熱延鋼板の製造方法。     Heating the steel material, subjecting it to hot rolling consisting of rough rolling and finish rolling, after finishing rolling, cooling, winding, and hot rolling steel sheet,
    The steel material in mass%,
    C: 0.06% to 0.1%,
    Si: 0.09% or less,
    Mn: 0.7% or more and 1.3% or less,
    P: 0.03% or less,
    S: 0.01% or less,
    Al: 0.1% or less,
    N: 0.01% or less,
    Ti: 0.14% to 0.20%
    V: 0.07% or more and 0.14% or less, with the balance being Fe and inevitable impurities,
    The heating temperature of the steel material is 1100 ° C. or more and 1350 ° C. or less, the finish rolling temperature of the finish rolling is 850 ° C. or more, the cooling is started within 3 seconds after finishing rolling, and the average cooling rate of the cooling is set. A method for producing a high-strength hot-rolled steel sheet at 20 ° C / s or higher, and a winding temperature of the winding is 550 ° C or higher and 700 ° C or lower.
  9. 　前記組成に加えてさらに、質量％でNb：0.01％以上0.05％以下を含有する請求項８に記載の高強度熱延鋼板の製造方法。 The method for producing a high-strength hot-rolled steel sheet according to claim 8, further comprising Nb: 0.01% or more and 0.05% or less by mass% in addition to the composition.
  10. 　前記組成に加えてさらに、Mo、W、Zr、Hfのいずれか１種以上を含有し、これらの含有量を、Mo：0.05％以下、W：0.05％以下、Zr：0.05％以下、Hf：0.05％以下に制限してなる請求項８又は９に記載の高強度熱延鋼板の製造方法。 In addition to the above composition, the composition further contains at least one of Mo, W, Zr, and Hf. The contents of Mo, 0.05% or less, W: 0.05% or less, Zr: 0.05% or less, Hf: The method for producing a high-strength hot-rolled steel sheet according to claim 8 or 9, which is limited to 0.05% or less.
  11. 　前記組成に加えてさらに、質量％で、O（酸素）、Se、Te、Po、As、Bi、Ge、Pb、Ga、In、Tl、Zn、Cd、Hg、Ag、Au、Pd、Pt、Co、Rh、Ir、Ru、Os、Tc、Re、Ta、Be、Sr、REM、B、Ni、Cr、Sb、Cu、Sn、Mg及びCaのうちの１種以上を合計で0.2％以下含有する請求項８～１０のいずれかに記載の高強度熱延鋼板の製造方法。 In addition to the above composition, in mass%, O (oxygen), Se, Te, Po, As, Bi, Ge, Pb, Ga, In, Tl, Zn, Cd, Hg, Ag, Au, Pd, Pt, Contains a total of 0.2% or less of one or more of Co, Rh, Ir, Ru, Os, Tc, Re, Ta, Be, Sr, REM, B, Ni, Cr, Sb, Cu, Sn, Mg and Ca The method for producing a high-strength hot-rolled steel sheet according to any one of claims 8 to 10.
  12. 　前記熱延鋼板の表面にめっき層を形成する請求項８～１１のいずれかに記載の高強度熱延鋼板の製造方法。 The method for producing a high-strength hot-rolled steel sheet according to any one of claims 8 to 11, wherein a plating layer is formed on the surface of the hot-rolled steel sheet.
  13. 　前記めっき層が亜鉛めっき層である請求項１２に記載の高強度熱延鋼板の製造方法。 The method for producing a high-strength hot-rolled steel sheet according to claim 12, wherein the plated layer is a galvanized layer.
  14. 　前記めっき層が合金化亜鉛めっき層である請求項１２に記載の高強度熱延鋼板の製造方法。
          The method for producing a high-strength hot-rolled steel sheet according to claim 12, wherein the plated layer is an alloyed galvanized layer.
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JPWO2013111556A1 (en) 2015-05-11
EP2808413A4 (en) 2016-05-04
EP2808413B1 (en) 2017-04-26
EP2808413A1 (en) 2014-12-03
US20150030880A1 (en) 2015-01-29
CN104053806A (en) 2014-09-17
JP5565534B2 (en) 2014-08-06
KR20140103340A (en) 2014-08-26

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