WO2013160938A1 - High strength cold-rolled steel plate of excellent ductility and manufacturing method therefor - Google Patents

High strength cold-rolled steel plate of excellent ductility and manufacturing method therefor Download PDF

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Publication number
WO2013160938A1
WO2013160938A1 PCT/JP2012/002807 JP2012002807W WO2013160938A1 WO 2013160938 A1 WO2013160938 A1 WO 2013160938A1 JP 2012002807 W JP2012002807 W JP 2012002807W WO 2013160938 A1 WO2013160938 A1 WO 2013160938A1
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Prior art keywords
phase
annealing
martensite
strength
volume fraction
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PCT/JP2012/002807
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French (fr)
Japanese (ja)
Inventor
英尚 川邉
横田　毅
瀬戸　一洋
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Ｊｆｅスチール株式会社
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Priority to PCT/JP2012/002807 priority Critical patent/WO2013160938A1/en
Publication of WO2013160938A1 publication Critical patent/WO2013160938A1/en

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Classifications

    • 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
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0447—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the heat treatment
    • C21D8/0473—Final recrystallisation annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/18—Hardening; Quenching with or without subsequent tempering
    • C21D1/25—Hardening, combined with annealing between 300 degrees Celsius and 600 degrees Celsius, i.e. heat refining ("Vergüten")
    • 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
    • C21D1/00—General methods or devices for heat treatment, e.g. annealing, hardening, quenching or tempering
    • C21D1/26—Methods of annealing
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • 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
    • C22C38/02—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
    • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • 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

Definitions

  • the present invention relates to a high-strength cold-rolled steel sheet suitable for use in automobile parts and the like that are required to be press-formed into a strict shape, and a method for producing the same, and in particular, actively uses expensive elements such as Cu, Ni, Cr, and Mo.
  • TS tensile strength
  • high-strength steel sheets of TS: 1180 MPa class or higher were often applied to light-worked parts, but recently, application to press parts with complex shapes has been studied.
  • steel sheets tend to have lower workability as the strength increases. Further, when the strength is increased to TS: 1180 MPa class or higher, the amount of additive elements such as C and Mn increases, and the weldability may be significantly reduced. This tendency is particularly strong in C. On the other hand, if it is difficult to contain a large amount of C, Si, or Mn from the viewpoints of weldability and chemical conversion properties, extremely expensive rare elements such as Nb, Cu, Ni, Cr, and Mo are required from the viewpoint of securing strength. May be positively added. Therefore, it is required to achieve both strength and formability with an alloy component that is low in content from the viewpoint of weldability and inexpensive from the economical aspect.
  • Patent Documents 1 to 4 disclose that high-strength cold-rolled steel is utilized by utilizing retained austenite by limiting steel components and structures, optimizing hot-rolling conditions, and annealing conditions. A technique for obtaining a rolled steel sheet is disclosed.
  • Patent Document 1 is mainly composed of tempered martensite obtained by performing tempering treatment after cooling to room temperature once in the annealing process, so that the volume fraction of retained austenite is small and sufficient ductility is achieved. Therefore, there is a problem that it cannot be applied to strict molding.
  • the base metal structure is mainly tempered martensite or bainite in which many carbides and dislocations exist in grains that are disadvantageous for ductility, the volume fraction of retained austenite is low. Since there are many, the component from which the outstanding ductility is obtained is disclosed. However, the content of C which is disadvantageous to weldability is large, and there remains a problem in that it is necessary to contain a large amount of expensive Cu and Ni.
  • Patent Document 3 has a problem in chemical conversion treatment property and post-coating corrosion resistance because the Cr content is essential, and there is a low volume fraction of retained austenite, which is a high dislocation density disadvantageous for ductility. There was a problem that excellent ductility could not be obtained due to the large volume fraction of nittic ferrite and martensite.
  • the technique described in Patent Document 4 needs to contain a large amount of C, which is disadvantageous for weldability, and has a high volume fraction of bainitic ferrite and martensite, which is disadvantageous for ductility. Therefore, excellent ductility is obtained. There was no problem.
  • the present invention advantageously solves the above-mentioned problems, reduces the content of C and Al, which are undesirable for weldability, and positively adds expensive elements such as Nb, Cu, Ni, Cr, and Mo.
  • the object is to provide a cold-rolled steel sheet together with its advantageous production method.
  • the gist configuration of the present invention is as follows. (1) In mass%, C: 0.16-0.20% Si: 1.0-2.0% Mn: 2.5-3.5% P: 0.030% or less, S: 0.0050% or less, Al: 0.005-0.1%, N: 0.01% or less, Ti: 0.001 to 0.050% and B: 0.0001 to 0.0050% And the balance has a component composition consisting of Fe and inevitable impurities, with a volume fraction, Ferrite phase: 40-65%, Martensite phase: 30-55% and residual austenite phase: 5-15% With a structure satisfying 0.5 to 5.0 martensite phases per unit area: 1 ⁇ m 2 in the cross section in the rolling direction, and having a tensile strength of 1180 MPa or more, high strength with excellent ductility Cold rolled steel sheet.
  • a high-strength cold-rolled steel sheet having excellent ductility and a tensile strength of 1180 MPa or more can be obtained without containing an expensive alloy element.
  • the high-strength cold-rolled steel sheet obtained by the present invention is suitable as an automobile part that is press-formed into a particularly severe shape.
  • the present invention will be specifically described below.
  • the inventors have obtained a volume fraction in a component system that is low C and does not contain Nb, Cu, Ni, Cr, or Mo. It has been found that the ductility is remarkably improved by forming a structure containing 40 to 65% ferrite phase, 30 to 55% martensite phase and 5 to 15% residual austenite phase.
  • the unit of the element content in the steel sheet is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.
  • C 0.16-0.20%
  • C is an element indispensable for forming a low-temperature transformation phase that contributes to securing the strength. If the amount of C is less than 0.16%, it is difficult to secure a desired steel sheet, and a desired amount of retained austenite cannot be obtained.
  • the C content exceeds 0.20%, not only the spot weldability is remarkably deteriorated, but also the low temperature transformation phase is excessively hardened, resulting in a decrease in formability. Therefore, the C content is in the range of 0.16 to 0.20%.
  • Si 1.0-2.0%
  • Si is an element that can be increased in strength without deteriorating elongation. Moreover, it has the effect
  • the Si content exceeds 2.0%, austenite formation is inhibited. Therefore, the Si content is in the range of 1.0 to 2.0%. The range is preferably 1.1 to 1.8%, more preferably 1.2 to 1.6%.
  • Mn 2.5-3.5%
  • Mn is an austenite stabilizing element and an essential element for obtaining a predetermined amount of retained austenite.
  • the content of 2.5% or more is necessary, but if it exceeds 3.5%, slab cracking occurs. Therefore, the Mn content is in the range of 2.5 to 3.5%. Preferably it is 2.6 to 3.0% of range.
  • P 0.030% or less P is preferably reduced as much as possible because P promotes grain boundary fracture by grain boundary segregation and adversely affects spot weldability, but 0.030% is acceptable. However, excessively reducing the amount of P lowers the production efficiency in the steel making process and increases the cost, so the lower limit of the amount of P is preferably about 0.001%.
  • S 0.0050% or less
  • S is a sulfide-based inclusion such as MnS in steel and is a starting point of cracking during stretch flange forming, which deteriorates formability. Therefore, it is preferable to reduce S as much as possible. Up to 0.0050% is acceptable. Preferably it is 0.0030% or less. However, excessive reduction of the amount of S is industrially difficult, and increases the desulfurization cost in the steel making process, so the lower limit of the amount of S is preferably about 0.0001%.
  • Al 0.005-0.1%
  • Al is used as a deoxidizer.
  • the Al amount needs to be 0.005% or more.
  • the Al amount exceeds 0.1%, the formability deteriorates due to an increase in inclusions such as alumina. Therefore, the Al content is in the range of 0.005 to 0.1%. Preferably it is 0.02 to 0.06% of range.
  • N 0.01% or less N combines with B to form BN, consumes B, and lowers the hardenability of solute B. Moreover, since it exists as an impurity element in ferrite and lowers the ductility by strain aging, it is preferable that the content is small, but an N content of up to 0.01% is acceptable. However, excessive reduction of the amount of N causes an increase in denitrification costs in the steelmaking process, so the lower limit of the amount of N is preferably about 0.0001%. More preferably, it is in the range of 0.0010 to 0.0050%.
  • Ti 0.001 to 0.050%
  • Ti is an element necessary for suppressing the formation of BN and expressing the hardenability by B by strongly fixing N as TiN.
  • Ti combines with C and N in steel to form fine carbides and nitrides, thereby suppressing the coarsening of crystal grains during heating, and the hot rolled sheet structure and the steel sheet structure after annealing Contributes effectively to uniform fine grains.
  • 0.001% or more of Ti is required.
  • the Ti content exceeds 0.050%, these effects tend to saturate. Ti precipitates are generated excessively, reducing the ductility of the ferrite phase, further hardening the hot-rolled sheet, and increasing the rolling load during hot rolling and cold rolling. Therefore, the Ti content is in the range of 0.001 to 0.050%. Preferably it is 0.005 to 0.025% of range.
  • B 0.0001-0.0050%
  • B is an element effective for enhancing hardenability, securing martensite, and achieving high strength.
  • the B amount needs to be 0.0001% or more.
  • the amount of B exceeds 0.0050%, the above effect is saturated. Therefore, the B amount is in the range of 0.0001 to 0.0050%. Preferably it is 0.0005 to 0.0020% of range.
  • components other than those described above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.
  • Ferrite phase volume fraction 40-65%
  • the ferrite phase is softer and more ductile than the hard martensite phase with high dislocation density, which is a low-temperature transformation phase from austenite, the bainite phase with carbides precipitated in the grains, and the bainitic ferrite with high dislocation density. It contributes to the improvement. In order to obtain this effect, it is necessary to contain 40% or more of a ferrite phase. On the other hand, if the ferrite phase is present in a volume fraction exceeding 65%, it is difficult to ensure a tensile strength of 1180 MPa or more.
  • the ferrite phase is in the range of 40 to 65% in volume fraction.
  • Martensite volume fraction 30-55%
  • the martensite phase contributes to the improvement of strength.
  • the volume fraction of the martensite phase exceeds 55%, it is preferable from the viewpoint of securing the strength, but it is difficult to secure a desired amount of ferrite phase and residual austenite phase that contribute to ductility. Therefore, the volume fraction of the martensite phase is in the range of 30 to 55%.
  • volume fraction of retained austenite phase 5-15%
  • the retained austenite phase is a structure that contributes to improvement of ductility by strain-induced transformation.
  • it is necessary to contain a residual austenite phase with a volume fraction of 5% or more.
  • the volume fraction of the retained austenite phase is 5 to 15%.
  • the number of martensite phases is less than 0.5 / ⁇ m 2 , there are few martensite phases and the tensile strength (TS) is insufficient, or coarse martensite phases are connected and present in high TS. The El will be lowered.
  • the martensite phase exists in the vicinity and the martensite phase surrounds the ferrite phase, so that the ferrite phase deforms and contributes to ductility. It becomes difficult. Accordingly, the number of martensite phases is in the range of 0.5 to 5.0 / ⁇ m 2 .
  • a steel slab having the above composition is annealed at 800 to 950 ° C, cooled to a cooling stop temperature of 200 to 500 ° C, then reheated to 750 to 850 ° C, and then an average cooling rate of 5 to 50 ° C / sec. Then, cool to the cooling stop temperature range of 350-450 ° C and let it stay in this temperature range for 100-1000 seconds.
  • the high-strength cold-rolled steel sheet which is the object of the present invention is obtained by such a production method, the obtained steel sheet may be subjected to skin pass rolling.
  • the process before hot finish rolling may be performed according to a conventional method.
  • a steel slab obtained by melting and casting steel prepared in the above component composition range can be used.
  • a steel slab obtained by melting and casting steel prepared in the above component composition range can be used.
  • not only continuous casting slabs and ingot-bundling slabs, but also thin slabs with a thickness of about 50 to 100 mm can be used.
  • direct heating without reheating is possible. It can use for a hot rolling process.
  • Hot rolling and cold rolling are not particularly limited, and may be performed according to a conventionally known method.
  • hot rolling may be performed at a hot rolling temperature of 850 to 950 ° C., wound at 450 to 650 ° C., and then cold rolled at a rolling reduction of 30 to 60%.
  • annealing is performed.
  • this annealing step is important, and the annealing is performed in two stages.
  • Annealing temperature 800-950 ° C
  • the annealing temperature in the first annealing is lower than 800 ° C.
  • the volume fraction of the ferrite phase increases during annealing, and the structure obtained after the first annealing becomes a structure mainly composed of a soft ferrite phase.
  • the enrichment of C in the austenite is not promoted, and a predetermined martensite phase and residual austenite cannot be obtained. As a result, it becomes difficult to secure a tensile strength of 1180 MPa or more.
  • the annealing temperature in the first annealing is in the range of 800 to 950 ° C. More preferably, it is in the range of 820 to 900 ° C.
  • Cooling stop temperature 200-500 ° C
  • the structure obtained after the first annealing becomes a structure containing mainly a bainite phase and a small amount of ferrite phase, and the formation of the ferrite phase is suppressed during the second annealing.
  • the volume fraction of the site phase becomes excessive and the TS increases, the ductility (El) decreases.
  • the cooling stop temperature after the first annealing is lower than 200 ° C.
  • the structure obtained after the first annealing is a structure mainly composed of martensite phase, and the formation of the ferrite phase is also suppressed during the second annealing.
  • the cooling stop temperature of the first annealing is set to 200 to 500 ° C. Also, if there is a lot of retained austenite after the first annealing, the second annealing will further promote the concentration of C in the austenite and promote the formation of retained austenite, so the cooling of the first annealing is stopped.
  • the temperature is preferably 350 to 450 ° C. at which the formation of retained austenite is promoted.
  • the cooling rate to the cooling stop temperature is not particularly limited, but is preferably about 10 to 50 ° C./second. Further, after the cooling is stopped, reheating may be continued, or after cooling to room temperature by cooling (air cooling), reheating may be performed.
  • Re-annealing temperature 750-850 ° C
  • the annealing temperature in the second annealing is set in the range of 750 to 850 ° C. More preferably, it is in the range of 770 to 830 ° C.
  • Average cooling rate 5 to 50 ° C./second
  • the cooling rate after the second annealing is important for obtaining a low temperature transformation phase having a desired volume fraction, and if the average cooling rate is lower than 5 ° C./second, The volume fraction of the ferrite phase generated in the process becomes too large, and it becomes difficult to secure a tensile strength of 1180 MPa or more.
  • the average cooling rate exceeds 50 ° C./second, the formation of the ferrite phase during cooling is suppressed, and the volume fraction of the martensite phase, which is a low-temperature transformation phase, increases from the austenite phase during annealing. It is easy to secure the 1180 MPa class, but the ductility decreases.
  • the average cooling rate after the second annealing is in the range of 5 to 50 ° C./second. More preferably, it is in the range of 10 to 35 ° C./second.
  • the cooling in this case is preferably gas cooling, but other methods such as furnace cooling, mist cooling, roll cooling, and water cooling can be used, or a combination thereof can also be used. .
  • Cooling stop temperature 350-450 ° C Since the low temperature transformation phase from austenite becomes harder as the transformation temperature is lower, the steel plate residence temperature after cooling is stopped is important for controlling the strength of the low temperature transformation phase.
  • the cooling stop temperature after the second annealing is lower than 350 ° C, it is cooled to a low temperature, so the low-temperature transformation phase is mainly a hard martensite phase, and it is easy to secure TS: 1180 MPa class, but the ductility is descend.
  • the cooling stop temperature after the second annealing is higher than 450 ° C., although the retained austenite phase is generated, the formation of the ferrite phase does not proceed, so that it becomes difficult to obtain excellent ductility.
  • the volume fraction of the martensite phase and the ferrite phase is controlled to ensure a tensile strength of 1180 MPa or more, while ensuring the desired volume fraction of the retained austenite phase and excellent ductility.
  • the cooling stop temperature after the second annealing must be in the range of 350 to 450 ° C.
  • Residence time at 350 to 450 ° C 100 to 1000 seconds Cooling stop temperature after the second annealing: If the residence time at 350 to 450 ° C is less than 100 seconds, it is difficult to obtain the desired retained austenite volume fraction In the cooling process to room temperature after residence, the untransformed austenite phase becomes the martensite phase, and the volume fraction of the martensite phase becomes excessive. As a result, the strength is increased, but the ductility is lowered. On the other hand, when retained for more than 1000 seconds, the formation of a retained austenite phase proceeds, so that ductility is improved, but it becomes difficult to obtain a tensile strength of 1180 MPa or more. Therefore, in order to secure a tensile strength of 1180 MPa or more and to obtain excellent ductility, the residence time at 350 to 450 ° C. needs to be in the range of 100 to 1000 seconds. The range is preferably 200 to 800 seconds.
  • Examples of the means for holding the steel sheet after stopping the cooling in the residence temperature range include a means for adjusting the temperature of the steel sheet to the residence temperature by providing a heat retaining device or the like in the downstream process of the cooling equipment after annealing. .
  • the steel plate after a residence is cooled to desired temperature by the conventionally well-known arbitrary methods.
  • the cold-rolled steel sheet finally obtained may be subjected to temper rolling (skin pass rolling) for the purpose of shape correction or surface roughness adjustment. Therefore, the crystal grains are expanded to form a rolled structure, and the ductility may be reduced. Therefore, the rolling reduction of skin pass rolling is preferably about 0.05% to 0.5%.
  • the volume fraction of the ferrite phase is the occupation of the phase existing in a square area of 10 ⁇ m ⁇ 10 ⁇ m square arbitrarily set by image analysis using a cross-sectional structure photograph of 1000 times magnification and cross-sectional SEM photographs of 1000 times and 3000 times The area was determined and used as the volume fraction of the ferrite phase.
  • the amount of retained austenite was determined by X-ray diffraction using Mo K ⁇ rays.
  • the volume fraction of the retained austenite phase was calculated.
  • the volume fraction of each phase is first distinguished from the ferrite phase and the low temperature transformation phase, first the volume fraction of the ferrite phase is determined, then the volume fraction of the residual austenite phase is determined by X-rays, and the remaining volume The fraction was judged as the martensite phase.
  • the number of martensite phases was determined by using a cross-sectional SEM photograph of 3000 times the number of martensite phases present in an arbitrarily set 10 ⁇ m ⁇ 10 ⁇ m square area, and this was taken as the number of martensite phases. .
  • the volume of each of the ferrite phase, martensite phase and residual austenite phase is reduced without reducing the amount of C in the steel sheet and actively including expensive elements such as Nb, Cu, Ni, Cr, Mo and V.
  • the fraction it is possible to obtain a high-strength cold-rolled steel sheet that is inexpensive and has excellent ductility and has a tensile strength (TS) of 1180 MPa or more.
  • TS tensile strength
  • the high-strength cold-rolled steel sheet of the present invention is also suitable for applications that require strict dimensional accuracy and workability, such as in the field of architecture and home appliances, in addition to automobile parts.

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  • Engineering & Computer Science (AREA)
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Abstract

A high strength cold-rolled steel plate with tensile strength of 1180 MPa or more and excellent ductility is provided as a result of the steel plate: having a component composition containing, in mass%, C: 0.16 - 0.20%, Si: 1.0 - 2.0%, Mn: 2.5 - 3.5%, P: 0.030% or less, S: 0.0050% or less, Al: 0.005 - 0.1%, N: 0.01% or less, Ti: 0.001 - 0.050%, and B: 0.0001 - 0.0050%, with the remainder being obtained from Fe and unavoidable impurities; comprising, in volume percent, 40 - 65% ferrite phase, 30 - 55% martensite phase, and 5 - 15% residual austenite phase; and having a structure in which the number of martensite phases per 1 µm2 unit area in a cross-section in the rolling direction is 0.5 - 5.0.

Description

延性に優れる高強度冷延鋼板およびその製造方法High strength cold-rolled steel sheet having excellent ductility and method for producing the same
　本発明は、厳しい形状にプレス成形されることが要求される自動車部品などに供して好適な高強度冷延鋼板およびその製造方法に関し、特にCuやNi，Cr，Moなど高価な元素を積極的に含有させることなしに、残留オーステナイトを活用し、また金属組織をフェライト相とマルテンサイト相を主体とした均一な組織とすることにより、延性の向上と共に、引張強度（TS）：1180MPa以上という高強度を併せて実現しようとするものである。 TECHNICAL FIELD The present invention relates to a high-strength cold-rolled steel sheet suitable for use in automobile parts and the like that are required to be press-formed into a strict shape, and a method for producing the same, and in particular, actively uses expensive elements such as Cu, Ni, Cr, and Mo. By using retained austenite without inclusion in the steel and making the metal structure a uniform structure mainly composed of ferrite and martensite phases, the ductility is improved and the tensile strength (TS) is as high as 1180 MPa or more. It tries to realize strength together.
　近年、衝突安全性の向上および車体軽量化による燃費向上の観点から、自動車車体に対して高強度鋼板の適用が拡大しつつあり、成形性に優れる高強度鋼板に対するニーズが高まっている。従来、ＴＳ：1180MPa級以上の高強度鋼板は、軽加工部品に適用されることが多かったが、最近では、複雑形状のプレス部品への適用が検討されている。 In recent years, the application of high-strength steel sheets to automobile bodies has been expanding from the viewpoint of improving collision safety and fuel efficiency by reducing the weight of the car body, and there is an increasing need for high-strength steel sheets having excellent formability. Conventionally, high-strength steel sheets of TS: 1180 MPa class or higher were often applied to light-worked parts, but recently, application to press parts with complex shapes has been studied.
　しかしながら、鋼板は、一般に、高強度化に伴い加工性が低下する傾向にある。また、ＴＳ：1180MPa級以上に高強度化する場合、ＣやMnなどの添加元素量が増加し、溶接性が著しく低下する場合がある。特にＣはこの傾向が強い。一方、溶接性や化成処理性の観点から、ＣやSi，Mnを多量に含有することが困難な場合、強度確保の観点から、Nb，Cu，Ni，Cr，Moなどの極めて高価な希少元素を積極的に添加する場合がある。
　従って、溶接性の点から含有量が少なく、かつ経済性の面から安価な合金成分で、強度と成形性を両立させることが要求されている。 However, generally, steel sheets tend to have lower workability as the strength increases. Further, when the strength is increased to TS: 1180 MPa class or higher, the amount of additive elements such as C and Mn increases, and the weldability may be significantly reduced. This tendency is particularly strong in C. On the other hand, if it is difficult to contain a large amount of C, Si, or Mn from the viewpoints of weldability and chemical conversion properties, extremely expensive rare elements such as Nb, Cu, Ni, Cr, and Mo are required from the viewpoint of securing strength. May be positively added.
Therefore, it is required to achieve both strength and formability with an alloy component that is low in content from the viewpoint of weldability and inexpensive from the economical aspect.
　成形性に優れた高強度冷延鋼板に関する従来技術として、例えば特許文献１～４に、鋼成分や組織の限定、熱延条件、焼鈍条件の最適化により、残留オーステナイトを活用して高強度冷延鋼板を得る技術が開示されている。 As conventional technologies related to high strength cold-rolled steel sheets with excellent formability, for example, Patent Documents 1 to 4 disclose that high-strength cold-rolled steel is utilized by utilizing retained austenite by limiting steel components and structures, optimizing hot-rolling conditions, and annealing conditions. A technique for obtaining a rolled steel sheet is disclosed.
特開２００４－２３８６７９号公報JP 2004-238679 A 特開２００５－１７９７０３号公報JP 2005-179703 A 特開２００７－１９７８１９号公報JP 2007-197819 A 特開２００８－１２７５８１号公報JP 2008-127581 A
　しかしながら、特許文献１に記載の技術は、焼鈍過程において一旦室温まで冷却してから焼戻し処理を施して得られる焼戻しマルテンサイトが主体であることから、残留オーステナイトの体積分率が少なく、十分な延性が得られないため、厳しい成形には適用できないという問題があった。
　特許文献２に記載の技術は、ベースとなる母相の金属組織が、延性に不利な粒内に炭化物や転位が多く存在する焼戻しマルテンサイトまたはベイナイト主体であるものの、残留オーステナイトの体積分率が多いため、優れた延性が得られる成分が開示されている。しかしながら、溶接性に不利なＣの含有量が多く、また高価なCuやNiを多量に含有させる必要があるところに問題を残していた。
　特許文献３に記載の技術は、Crの含有が必須であることから、化成処理性および塗装後耐食性に問題があり、また残留オーステナイトの体積分率が少なく、延性に不利な転位密度の高いベイニティックフェライトとマルテンサイトの体積分率が多いため、優れた延性が得られないという問題があった。
　特許文献４に記載の技術は、溶接性に不利なＣを多量に含有させる必要があり、また延性に不利なベイニティックフェライトとマルテンサイトの体積分率が多いため、優れた延性が得られないという問題があった。 However, the technique described in Patent Document 1 is mainly composed of tempered martensite obtained by performing tempering treatment after cooling to room temperature once in the annealing process, so that the volume fraction of retained austenite is small and sufficient ductility is achieved. Therefore, there is a problem that it cannot be applied to strict molding.
In the technique described in Patent Document 2, although the base metal structure is mainly tempered martensite or bainite in which many carbides and dislocations exist in grains that are disadvantageous for ductility, the volume fraction of retained austenite is low. Since there are many, the component from which the outstanding ductility is obtained is disclosed. However, the content of C which is disadvantageous to weldability is large, and there remains a problem in that it is necessary to contain a large amount of expensive Cu and Ni.
The technique described in Patent Document 3 has a problem in chemical conversion treatment property and post-coating corrosion resistance because the Cr content is essential, and there is a low volume fraction of retained austenite, which is a high dislocation density disadvantageous for ductility. There was a problem that excellent ductility could not be obtained due to the large volume fraction of nittic ferrite and martensite.
The technique described in Patent Document 4 needs to contain a large amount of C, which is disadvantageous for weldability, and has a high volume fraction of bainitic ferrite and martensite, which is disadvantageous for ductility. Therefore, excellent ductility is obtained. There was no problem.
　本発明は、上記の問題を有利に解決するもので、溶接性に好ましくないＣやAlの含有量を少なくし、しかも高価な元素であるNb，Cu、Ni、Cr、Mo等を積極的に含有させない成分系で、残留オーステナイト相の体積分率を厳密に制御すると共に、フェライト相とマルテンサイト相主体の組織とすることにより、延性に優れかつ引張強度（TS）が1180MPa以上である高強度冷延鋼板を、その有利な製造方法と共に提供することを目的とする。 The present invention advantageously solves the above-mentioned problems, reduces the content of C and Al, which are undesirable for weldability, and positively adds expensive elements such as Nb, Cu, Ni, Cr, and Mo. High strength with excellent ductility and tensile strength (TS) of 1180 MPa or more by strictly controlling the volume fraction of the retained austenite phase and making the structure mainly composed of ferrite and martensite phases. The object is to provide a cold-rolled steel sheet together with its advantageous production method.
　すなわち、本発明の要旨構成は以下のとおりである。
（１）質量％で、
　　Ｃ：0.16～0.20％、
　　Si：1.0～2.0％、
　　Mn：2.5～3.5％、
　　Ｐ：0.030％以下、
　　Ｓ：0.0050％以下、
　　Al：0.005～0.1％、
　　Ｎ：0.01％以下、
　　Ti：0.001～0.050％および
　　Ｂ：0.0001～0.0050％
を含有し、残部がFe及び不可避的不純物からなる成分組成を有し、体積分率で、
　　フェライト相：40～65％、
　　マルテンサイト相：30～55％および
　　残留オーステナイト相：５～15％
を含み、圧延方向断面において単位面積：１μm2当たりのマルテンサイト相の数が 0.5～5.0個を満足する組織を有し、引張強度：1180MPa以上であることを特徴とする、延性に優れる高強度冷延鋼板。 That is, the gist configuration of the present invention is as follows.
(1) In mass%,
C: 0.16-0.20%
Si: 1.0-2.0%
Mn: 2.5-3.5%
P: 0.030% or less,
S: 0.0050% or less,
Al: 0.005-0.1%,
N: 0.01% or less,
Ti: 0.001 to 0.050% and B: 0.0001 to 0.0050%
And the balance has a component composition consisting of Fe and inevitable impurities, with a volume fraction,
Ferrite phase: 40-65%,
Martensite phase: 30-55% and residual austenite phase: 5-15%
With a structure satisfying 0.5 to 5.0 martensite phases per unit area: 1 μm 2 in the cross section in the rolling direction, and having a tensile strength of 1180 MPa or more, high strength with excellent ductility Cold rolled steel sheet.
（２）上記（１）に記載の成分組成からなる鋼スラブを、熱間圧延し、ついで酸洗後、冷間圧延したのち、焼鈍を施して高強度冷延鋼板を製造するに際し、
　上記焼鈍工程において、800～950℃で焼鈍後、冷却停止温度：200～500℃まで冷却し、ついで750～850℃に再加熱後、平均冷却速度：５～50℃/秒の速度で、350～450℃の冷却停止温度域まで冷却し、この温度域に100～1000秒滞留させることを特徴とする、延性に優れる高強度冷延鋼板の製造方法。 (2) When a steel slab having the composition described in (1) above is hot-rolled, then pickled, cold-rolled, and then annealed to produce a high-strength cold-rolled steel sheet,
In the above annealing process, after annealing at 800 to 950 ° C., the cooling stop temperature is cooled to 200 to 500 ° C., then reheated to 750 to 850 ° C., and the average cooling rate is 350 ° C. at a rate of 5 to 50 ° C./second. A method for producing a high-strength cold-rolled steel sheet with excellent ductility, characterized by cooling to a cooling stop temperature range of up to 450 ° C. and retaining in this temperature range for 100 to 1000 seconds.
　本発明によれば、高価な合金元素を含有させることなしに、延性に優れ、しかも引張強度が1180MPa以上の高強度冷延鋼板を得ることができる。そして、本発明により得られる高強度冷延鋼板は、特に厳しい形状にプレス成形される自動車部品として好適である。 According to the present invention, a high-strength cold-rolled steel sheet having excellent ductility and a tensile strength of 1180 MPa or more can be obtained without containing an expensive alloy element. The high-strength cold-rolled steel sheet obtained by the present invention is suitable as an automobile part that is press-formed into a particularly severe shape.
　以下、本発明を具体的に説明する。
　さて、発明者らは、高強度冷延鋼板の延性の向上に関し、鋭意検討を重ねた結果、低Ｃで、かつNb，Cu、Ni、Cr、Moを含有しない成分系において、体積分率で、40～65％のフェライト相、30～55％のマルテンサイト相、５～15％の残留オーステナイト相を含む組織とすることにより、延性の向上が顕著となることを見出した。
　以下、本発明の成分組成および組織の限定理由について具体的に説明する。なお、鋼板中の元素の含有量の単位は何れも「質量％」であるが、以下、特に断らない限り、単に「％」で示す。 The present invention will be specifically described below.
As a result of intensive investigations on the improvement of ductility of high-strength cold-rolled steel sheets, the inventors have obtained a volume fraction in a component system that is low C and does not contain Nb, Cu, Ni, Cr, or Mo. It has been found that the ductility is remarkably improved by forming a structure containing 40 to 65% ferrite phase, 30 to 55% martensite phase and 5 to 15% residual austenite phase.
Hereinafter, the reasons for limiting the component composition and structure of the present invention will be specifically described. The unit of the element content in the steel sheet is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.
　まず、本発明における鋼の成分組成の適正範囲およびその限定理由は以下のとおりである。
Ｃ：0.16～0.20％
　Ｃは、強度確保に寄与する低温変態相の形成に不可欠の元素である。Ｃ量が0.16％に満たないと、所望の鋼板を確保することが難しく、また所望量の残留オーステナイトが得られない。一方、Ｃ量が0.20％を超えると、スポット溶接性が著しく劣化するだけでなく、低温変態相が過度に硬質化して成形性の低下を招く。そのため、Ｃ量は0.16～0.20％の範囲とする。 First, the appropriate range of the component composition of steel in the present invention and the reasons for limitation are as follows.
C: 0.16-0.20%
C is an element indispensable for forming a low-temperature transformation phase that contributes to securing the strength. If the amount of C is less than 0.16%, it is difficult to secure a desired steel sheet, and a desired amount of retained austenite cannot be obtained. On the other hand, when the C content exceeds 0.20%, not only the spot weldability is remarkably deteriorated, but also the low temperature transformation phase is excessively hardened, resulting in a decrease in formability. Therefore, the C content is in the range of 0.16 to 0.20%.
Si：1.0～2.0％
　Siは、伸びを劣化させずに高強度化が可能な元素である。また、セメンタイトの析出を抑制し、残留オーステナイトを残存させやすくする作用もある。上記の効果を得るには、1.0％以上含有させる必要がある。一方、Si量が2.0％を超えると、オーステナイトの形成が阻害される。そのため、Si量は1.0～2.0％の範囲とする。好ましくは1.1～1.8％の範囲、より好ましくは1.2～1.6％の範囲である。 Si: 1.0-2.0%
Si is an element that can be increased in strength without deteriorating elongation. Moreover, it has the effect | action which suppresses precipitation of cementite and makes it easy to leave a retained austenite. In order to acquire said effect, it is necessary to contain 1.0% or more. On the other hand, when the Si content exceeds 2.0%, austenite formation is inhibited. Therefore, the Si content is in the range of 1.0 to 2.0%. The range is preferably 1.1 to 1.8%, more preferably 1.2 to 1.6%.
Mn：2.5～3.5％
　Mnは、オーステナイト安定化元素であり、所定量の残留オーステナイトを得るために必須の元素である。上記の作用を得るためには2.5％以上の含有が必要であるが、3.5％を超えて含有させるとスラブ割れが生じる。そのため、Mn量は2.5～3.5％の範囲とする。好ましくは2.6～3.0％の範囲である。 Mn: 2.5-3.5%
Mn is an austenite stabilizing element and an essential element for obtaining a predetermined amount of retained austenite. In order to obtain the above effect, the content of 2.5% or more is necessary, but if it exceeds 3.5%, slab cracking occurs. Therefore, the Mn content is in the range of 2.5 to 3.5%. Preferably it is 2.6 to 3.0% of range.
Ｐ：0.030％以下
　Ｐは、粒界偏析により粒界破壊を助長し、またスポット溶接性に悪影響を及ぼすため、極力低減することが好ましいが、0.030％までは許容できる。しかし、Ｐ量を過度に低減することは製鋼工程での生産能率が低下し、高コストとなるため、Ｐ量の下限は0.001％程度とすることが好ましい。 P: 0.030% or less P is preferably reduced as much as possible because P promotes grain boundary fracture by grain boundary segregation and adversely affects spot weldability, but 0.030% is acceptable. However, excessively reducing the amount of P lowers the production efficiency in the steel making process and increases the cost, so the lower limit of the amount of P is preferably about 0.001%.
Ｓ：0.0050％以下
　Ｓは、鋼中でMnＳなどの硫化物系介在物となり、伸びフランジ成形時の割れの起点となって成形性を劣化させるので、極力低減することが好ましいが、Ｓ量が0.0050％までは許容できる。好ましくは0.0030％以下である。しかし、Ｓ量の過度の低減は工業的に困難であり、製鋼工程における脱硫コストの増加を招くので、Ｓ量の下限は0.0001％程度とすることが好ましい。 S: 0.0050% or less S is a sulfide-based inclusion such as MnS in steel and is a starting point of cracking during stretch flange forming, which deteriorates formability. Therefore, it is preferable to reduce S as much as possible. Up to 0.0050% is acceptable. Preferably it is 0.0030% or less. However, excessive reduction of the amount of S is industrially difficult, and increases the desulfurization cost in the steel making process, so the lower limit of the amount of S is preferably about 0.0001%.
Al：0.005～0.1％
　Alは、脱酸剤として使用される。脱酸作用を得るためにはAl量を0.005％以上とすることが必要であるが、Al量が0.1％を超えると、アルミナなどの介在物増加による成形性の劣化が生じる。従って、Al量は0.005～0.1％の範囲とする。好ましくは0.02～0.06％の範囲である。 Al: 0.005-0.1%
Al is used as a deoxidizer. In order to obtain a deoxidizing action, the Al amount needs to be 0.005% or more. However, if the Al amount exceeds 0.1%, the formability deteriorates due to an increase in inclusions such as alumina. Therefore, the Al content is in the range of 0.005 to 0.1%. Preferably it is 0.02 to 0.06% of range.
Ｎ：0.01％以下
　Ｎは、Ｂと結合しＢＮを形成してＢを消費し、固溶Ｂによる焼入れ性を低下させる。また、フェライト中で不純物元素として存在し、ひずみ時効により延性を低下させるので、少ない方が好ましいが、Ｎ量が0.01％までは許容できる。しかし、Ｎ量の過度の低減は製鋼工程における脱窒コストの増加を招くので、Ｎ量の下限は0.0001％程度とすることが好ましい。より好ましくは0.0010～0.0050％の範囲である。 N: 0.01% or less N combines with B to form BN, consumes B, and lowers the hardenability of solute B. Moreover, since it exists as an impurity element in ferrite and lowers the ductility by strain aging, it is preferable that the content is small, but an N content of up to 0.01% is acceptable. However, excessive reduction of the amount of N causes an increase in denitrification costs in the steelmaking process, so the lower limit of the amount of N is preferably about 0.0001%. More preferably, it is in the range of 0.0010 to 0.0050%.
Ti：0.001～0.050％
　Tiは、ＮをTiＮとして強く固定することにより、ＢＮの形成を抑制し、Ｂによる焼入れ性を発現させるのに必要な元素である。また、Tiは、鋼中でＣやＮと結合して微細な炭化物や窒化物を形成することにより、加熱時における結晶粒の粗大化を抑制し、熱延板組織および焼鈍後の鋼板組織の細粒均一化に有効に寄与する。これらの効果を得るには0.001％以上のTi含有を必要とするが、Ti量が0.050％を超えるとこれらの効果は飽和する傾向にあり、またTiを過度に含有させると、フェライト相中にTiの析出物が過剰に生成し、フェライト相の延性を低下させ、さらには熱延板が硬質化し、熱間圧延時および冷間圧延時の圧延荷重が増大する。従って、Ti量は0.001～0.050％の範囲とする。好ましくは0.005～0.025％の範囲である。 Ti: 0.001 to 0.050%
Ti is an element necessary for suppressing the formation of BN and expressing the hardenability by B by strongly fixing N as TiN. In addition, Ti combines with C and N in steel to form fine carbides and nitrides, thereby suppressing the coarsening of crystal grains during heating, and the hot rolled sheet structure and the steel sheet structure after annealing Contributes effectively to uniform fine grains. In order to obtain these effects, 0.001% or more of Ti is required. However, when the Ti content exceeds 0.050%, these effects tend to saturate. Ti precipitates are generated excessively, reducing the ductility of the ferrite phase, further hardening the hot-rolled sheet, and increasing the rolling load during hot rolling and cold rolling. Therefore, the Ti content is in the range of 0.001 to 0.050%. Preferably it is 0.005 to 0.025% of range.
Ｂ：0.0001～0.0050％
　Ｂは、焼入れ性を高め、マルテンサイトを確保して高強度化を達成するのに有効な元素である。かかる効果を得るためには、Ｂ量を0.0001％以上とする必要がある。一方、Ｂ量が0.0050％を超えると上記効果は飽和する。従って、Ｂ量は0.0001～0.0050％の範囲とする。好ましくは0.0005～0.0020％の範囲である。
　なお、本発明の鋼板において、上記以外の成分はFeおよび不可避的不純物である。ただし、本発明の効果を損なわない範囲内であれば、上記以外の成分の含有を拒むものではない。 B: 0.0001-0.0050%
B is an element effective for enhancing hardenability, securing martensite, and achieving high strength. In order to obtain such an effect, the B amount needs to be 0.0001% or more. On the other hand, when the amount of B exceeds 0.0050%, the above effect is saturated. Therefore, the B amount is in the range of 0.0001 to 0.0050%. Preferably it is 0.0005 to 0.0020% of range.
In the steel sheet of the present invention, components other than those described above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.
　次に、本発明にとって重要な要件の一つである鋼の組織の適正範囲およびその限定理由について説明する。
フェライト相の体積分率：40～65％
　フェライト相は、オーステナイトからの低温変態相である高転位密度の硬質なマルテンサイト相や、粒内に炭化物が析出しているベイナイト相、転位密度の高いベイニティックフェライトよりも軟質であり、延性の向上に寄与する。この効果を得るためには、40％以上のフェライト相を含有させる必要がある。一方、フェライト相が体積分率で65％を超えて存在すると1180MPa以上の引張強度の確保が困難となる。また、フェライト相の体積分率が多くなると、フェライト相とマルテンサイト相を主体としつつ、延性に寄与する残留オーステナイト相を確保することが困難となる。従って、フェライト相は体積分率で40～65％の範囲とする。 Next, the appropriate range of the steel structure, which is one of the important requirements for the present invention, and the reason for the limitation will be described.
Ferrite phase volume fraction: 40-65%
The ferrite phase is softer and more ductile than the hard martensite phase with high dislocation density, which is a low-temperature transformation phase from austenite, the bainite phase with carbides precipitated in the grains, and the bainitic ferrite with high dislocation density. It contributes to the improvement. In order to obtain this effect, it is necessary to contain 40% or more of a ferrite phase. On the other hand, if the ferrite phase is present in a volume fraction exceeding 65%, it is difficult to ensure a tensile strength of 1180 MPa or more. Further, when the volume fraction of the ferrite phase increases, it becomes difficult to secure a retained austenite phase that contributes to ductility while mainly including the ferrite phase and the martensite phase. Therefore, the ferrite phase is in the range of 40 to 65% in volume fraction.
マルテンサイト相の体積分率：30～55％
　マルテンサイト相は強度の向上に寄与する。軟質なフェライト相の体積分率を確保しつつ所望の強度を確保するためには、30％以上のマルテンサイト相を含有させる必要がある。一方、マルテンサイト相の体積分率が55％を超えた場合、強度確保の面からは好ましいが、延性に寄与する所望量のフェライト相および残留オーステナイト相の確保が困難となる。従って、マルテンサイト相の体積分率は30～55％の範囲とする。 Martensite volume fraction: 30-55%
The martensite phase contributes to the improvement of strength. In order to ensure the desired strength while ensuring the volume fraction of the soft ferrite phase, it is necessary to contain 30% or more of the martensite phase. On the other hand, when the volume fraction of the martensite phase exceeds 55%, it is preferable from the viewpoint of securing the strength, but it is difficult to secure a desired amount of ferrite phase and residual austenite phase that contribute to ductility. Therefore, the volume fraction of the martensite phase is in the range of 30 to 55%.
残留オーステナイト相の体積分率：５～15％
　残留オーステナイト相は、歪誘起変態により延性の向上に寄与する組織である。特に高延性化するためには体積分率で５％以上の残留オーステナイト相を含有させることが必要である。一方、体積分率で15％を超えて含有させると、フェライト相とマルテンサイト相の体積分率をバランスさせて高い強度と優れた延性のバランスを確保することが困難となる。よって、残留オーステナイト相の体積分率は５～15％とする。 Volume fraction of retained austenite phase: 5-15%
The retained austenite phase is a structure that contributes to improvement of ductility by strain-induced transformation. In particular, in order to increase ductility, it is necessary to contain a residual austenite phase with a volume fraction of 5% or more. On the other hand, if the volume fraction exceeds 15%, it becomes difficult to balance the volume fraction of the ferrite phase and the martensite phase to ensure a balance between high strength and excellent ductility. Therefore, the volume fraction of the retained austenite phase is 5 to 15%.
圧延方向断面におけマルテンサイト相の数：0.5～5.0個／μm2
　マルテンサイト相が連結してフェライトを囲むように存在していると、フェライト相の延性への寄与は小さい。フェライト相が変形し、延性（Ｅｌ）に寄与するためには、フェライト相中にマルテンサイト相が孤立して存在していることが必要である。マルテンサイト相の存在状態は厳密には３次元で評価する必要があるが、鋼中の組織を３次元化して評価するのはコスト、時間がかかりすぎるので、簡便な２次元の圧延方向に平行な断面で、板厚1/4面で測定し、評価するものとする。
　この評価で、マルテンサイト相の数が0.5個／μm2未満の場合、マルテンサイト相が少なく、引張強度（ＴＳ）が不足するか、または粗大なマルテンサイト相が連結して存在し、高ＴＳ低Ｅｌ化することになる。一方、5.0個／μm2超の場合、マルテンサイト相が近接して存在することになり、フェライト相の周囲をマルテンサイト相が囲むようになるので、フェライト相が変形して延性に寄与することが困難となる。従って、マルテンサイト相の数は0.5～5.0個／μm2の範囲とする。 Number of martensite phases in the cross section in the rolling direction: 0.5 to 5.0 / μm 2
If the martensite phase is connected to surround the ferrite, the contribution of the ferrite phase to the ductility is small. In order for the ferrite phase to be deformed and contribute to ductility (El), it is necessary that the martensite phase is present in isolation in the ferrite phase. Strictly speaking, it is necessary to evaluate the presence of the martensite phase in three dimensions. However, it is too costly and time-consuming to evaluate the structure in steel in three dimensions, so it is parallel to the simple two-dimensional rolling direction. With a simple cross section, measurement shall be made with a quarter thickness of the plate and evaluated.
In this evaluation, when the number of martensite phases is less than 0.5 / μm 2 , there are few martensite phases and the tensile strength (TS) is insufficient, or coarse martensite phases are connected and present in high TS. The El will be lowered. On the other hand, in the case of more than 5.0 pieces / μm 2 , the martensite phase exists in the vicinity and the martensite phase surrounds the ferrite phase, so that the ferrite phase deforms and contributes to ductility. It becomes difficult. Accordingly, the number of martensite phases is in the range of 0.5 to 5.0 / μm 2 .
　次に本発明の高強度冷延鋼板の製造方法について説明する。
　上記の成分組成を有する鋼スラブを、800～950℃で焼鈍後、冷却停止温度：200～500℃まで冷却し、引き続き750～850℃に再加熱後、平均冷却速度：５～50℃／秒で、350～450℃の冷却停止温度域まで冷却し、この温度域に100～1000秒滞留させる。かかる製造方法により本発明の目的とする高強度冷延鋼板が得られるが、得られた鋼板にスキンパス圧延を施しても良い。 Next, the manufacturing method of the high-strength cold-rolled steel sheet of this invention is demonstrated.
A steel slab having the above composition is annealed at 800 to 950 ° C, cooled to a cooling stop temperature of 200 to 500 ° C, then reheated to 750 to 850 ° C, and then an average cooling rate of 5 to 50 ° C / sec. Then, cool to the cooling stop temperature range of 350-450 ° C and let it stay in this temperature range for 100-1000 seconds. Although the high-strength cold-rolled steel sheet which is the object of the present invention is obtained by such a production method, the obtained steel sheet may be subjected to skin pass rolling.
　以下、製造条件の適正範囲およびその限定理由について説明する。
　本発明において、熱間仕上げ圧延前の工程に関しては常法に従って行えばよく、例えば、上記の成分組成範囲に調製した鋼を溶製、鋳造して得られた鋼スラブを用いることができる。また、本発明においては、連続鋳造スラブ、造塊－分塊スラブは勿論のこと、厚み：50～100mm程度の薄スラブを用いることができ、特に薄スラブの場合は、再加熱なしに直接熱間圧延工程に供することができる。 Hereinafter, the appropriate range of manufacturing conditions and the reason for limitation will be described.
In the present invention, the process before hot finish rolling may be performed according to a conventional method. For example, a steel slab obtained by melting and casting steel prepared in the above component composition range can be used. In the present invention, not only continuous casting slabs and ingot-bundling slabs, but also thin slabs with a thickness of about 50 to 100 mm can be used. Particularly in the case of thin slabs, direct heating without reheating is possible. It can use for a hot rolling process.
　熱間圧延および冷間圧延についても特に制限はなく、従来公知の方法に従って行えばよい。例えば、熱間圧延を熱延出側温度：850～950℃で行い、450～650℃で巻き取ったのち、圧下率：30～60％で冷間圧延を行えばよい。 Hot rolling and cold rolling are not particularly limited, and may be performed according to a conventionally known method. For example, hot rolling may be performed at a hot rolling temperature of 850 to 950 ° C., wound at 450 to 650 ° C., and then cold rolled at a rolling reduction of 30 to 60%.
　ついで、焼鈍を施すが、本発明では、この焼鈍工程が重要であり、焼鈍を２段階に分けて行う。
焼鈍温度：800～950℃
　１回目の焼鈍における焼鈍温度が800℃より低い場合、焼鈍中にフェライト相の体積分率が多くなり、１回目の焼鈍後に得られる組織が軟質なフェライト相主体の組織となるため、２回目の焼鈍時にオーステナイト中へのＣの濃化が促進されず、所定のマルテンサイト相および残留オーステナイトが得られなくなる結果、1180 MPa以上の引張強度の確保が困難となる。一方、１回目における焼鈍が950℃を超えてオーステナイト単相の高温域まで加熱すると、オーステナイト粒径が過度に粗大化して、１回目の焼鈍後に得られる組織も粗大化し、また２回目の焼鈍時にフェライト相の生成が抑制されてマルテンサイト相の体積分率が過剰となるため、高ＴＳ化はするものの、延性（Ｅｌ）は低下する。従って、１回目の焼鈍における焼鈍温度は800～950℃の範囲とする。より好ましくは、820～900℃の範囲である。 Next, annealing is performed. In the present invention, this annealing step is important, and the annealing is performed in two stages.
Annealing temperature: 800-950 ° C
When the annealing temperature in the first annealing is lower than 800 ° C., the volume fraction of the ferrite phase increases during annealing, and the structure obtained after the first annealing becomes a structure mainly composed of a soft ferrite phase. During annealing, the enrichment of C in the austenite is not promoted, and a predetermined martensite phase and residual austenite cannot be obtained. As a result, it becomes difficult to secure a tensile strength of 1180 MPa or more. On the other hand, if the annealing at the first time exceeds 950 ° C. and is heated to the high temperature region of the austenite single phase, the austenite grain size becomes excessively coarse, and the structure obtained after the first annealing also becomes coarse, and also during the second annealing. Since the formation of the ferrite phase is suppressed and the volume fraction of the martensite phase becomes excessive, the ductility (El) is lowered although the TS is increased. Therefore, the annealing temperature in the first annealing is in the range of 800 to 950 ° C. More preferably, it is in the range of 820 to 900 ° C.
冷却停止温度：200～500℃
　１回目の焼鈍後の冷却停止温度が500℃を上回ると、１回目の焼鈍後に得られる組織はベイナイト相主体でフェライト相の少ない組織となり、２回目の焼鈍時にフェライト相の生成が抑制されてマルテンサイト相の体積分率が過剰となり、高ＴＳ化するものの、延性（Ｅｌ）が低下する。一方、１回目の焼鈍後の冷却停止温度が200℃を下回った場合、１回目の焼鈍後に得られる組織はマルテンサイト相主体の組織となり、やはり２回目の焼鈍時にフェライト相の生成が抑制されてマルテンサイト相の体積分率が過剰となり、高ＴＳ化はするものの、延性（Ｅｌ）は低下する。そのため、１回目の焼鈍の冷却停止温度は200～500℃とする。また、１回目の焼鈍後に残留オーステナイトが多く存在していると、２回目の焼鈍によってさらにオーステナイト中へのＣ濃化がすすみ、残留オーステナイトの生成が促進されるので、１回目の焼鈍の冷却停止温度は、残留オーステナイトの生成が促進される350～450℃とするのが好ましい。
　なお、この冷却停止温度までの冷却速度は特に限定されることはないが、10～50℃／秒程度とするのが好適である。
　また、冷却停止後は、引き続き再加熱を行ってもよいし、放冷（空冷）にて一旦室温まで冷却したのち、再加熱を行ってもよい。 Cooling stop temperature: 200-500 ° C
When the cooling stop temperature after the first annealing exceeds 500 ° C., the structure obtained after the first annealing becomes a structure containing mainly a bainite phase and a small amount of ferrite phase, and the formation of the ferrite phase is suppressed during the second annealing. Although the volume fraction of the site phase becomes excessive and the TS increases, the ductility (El) decreases. On the other hand, when the cooling stop temperature after the first annealing is lower than 200 ° C., the structure obtained after the first annealing is a structure mainly composed of martensite phase, and the formation of the ferrite phase is also suppressed during the second annealing. The volume fraction of the martensite phase becomes excessive and the TS increases, but the ductility (El) decreases. Therefore, the cooling stop temperature of the first annealing is set to 200 to 500 ° C. Also, if there is a lot of retained austenite after the first annealing, the second annealing will further promote the concentration of C in the austenite and promote the formation of retained austenite, so the cooling of the first annealing is stopped. The temperature is preferably 350 to 450 ° C. at which the formation of retained austenite is promoted.
The cooling rate to the cooling stop temperature is not particularly limited, but is preferably about 10 to 50 ° C./second.
Further, after the cooling is stopped, reheating may be continued, or after cooling to room temperature by cooling (air cooling), reheating may be performed.
再焼鈍温度（再加熱温度）：750～850℃
　２回目の焼鈍における焼鈍温度が750℃より低い場合、焼鈍中のオーステナイト相の体積分率が少なく、最終的に得られるマルテンサイト相の体積分率が少なくなるため、1180 MPa以上の引張強度の確保が困難となる。一方、２回目における焼鈍温度が850℃を超えると、焼鈍中にオーステナイト相が粗大化し、その後の冷却、保持中におけるフェライト相の生成が抑制されるため、高ＴＳ化するものの、延性（Ｅｌ）は低下する。従って、２回目の焼鈍における焼鈍温度は750～850℃の範囲とする。より好ましくは、770～830℃の範囲である。 Re-annealing temperature (re-heating temperature): 750-850 ° C
When the annealing temperature in the second annealing is lower than 750 ° C, the volume fraction of the austenite phase during annealing is small and the volume fraction of the finally obtained martensite phase is small. It becomes difficult to secure. On the other hand, if the annealing temperature in the second time exceeds 850 ° C., the austenite phase becomes coarse during annealing, and the formation of ferrite phase during subsequent cooling and holding is suppressed, so that although the TS increases, ductility (El) Will decline. Therefore, the annealing temperature in the second annealing is set in the range of 750 to 850 ° C. More preferably, it is in the range of 770 to 830 ° C.
平均冷却速度：５～50℃／秒
　２回目の焼鈍後の冷却速度は、所望の体積分率の低温変態相を得るために重要であり、平均冷却速度が５℃／秒より遅いと、冷却過程で生成するフェライト相の体積分率が多くなりすぎ、1180 MPa以上の引張強度の確保が困難となる。一方、平均冷却速度が50℃／秒を超えると、冷却中におけるフェライト相の生成が抑制され、焼鈍中のオーステナイト相から低温変態相であるマルテンサイト相の体積分率が増加するため、ＴＳ：1180 MPa級の確保は容易であるが、延性は低下する。それ故、２回目の焼鈍後の平均冷却速度は５～50℃／秒の範囲とする。より好ましくは10～35℃／秒の範囲である。
　なお、この場合の冷却は、ガス冷却とすることが好ましいが、その他、炉冷、ミスト冷却、ロール冷却および水冷などの方法を用いることができ、またはそれらを組み合わせて使用することも可能である。 Average cooling rate: 5 to 50 ° C./second The cooling rate after the second annealing is important for obtaining a low temperature transformation phase having a desired volume fraction, and if the average cooling rate is lower than 5 ° C./second, The volume fraction of the ferrite phase generated in the process becomes too large, and it becomes difficult to secure a tensile strength of 1180 MPa or more. On the other hand, when the average cooling rate exceeds 50 ° C./second, the formation of the ferrite phase during cooling is suppressed, and the volume fraction of the martensite phase, which is a low-temperature transformation phase, increases from the austenite phase during annealing. It is easy to secure the 1180 MPa class, but the ductility decreases. Therefore, the average cooling rate after the second annealing is in the range of 5 to 50 ° C./second. More preferably, it is in the range of 10 to 35 ° C./second.
The cooling in this case is preferably gas cooling, but other methods such as furnace cooling, mist cooling, roll cooling, and water cooling can be used, or a combination thereof can also be used. .
冷却停止温度：350～450℃
　オーステナイトからの低温変態相は、変態温度が低いほど硬くなるため、冷却停止後の鋼板滞留温度は、低温変態相の強度を制御するのに重要である。
　２回目の焼鈍後の冷却停止温度が350℃より低い場合、低温まで冷却されるため、低温変態相は硬質なマルテンサイト相主体となり、ＴＳ：1180 MPa級の確保は容易であるが、延性は低下する。一方、２回目の焼鈍後の冷却停止温度が450℃より高い場合、残留オーステナイト相は生成するもののフェライト相の生成が進行しないため、優れた延性を得ることが困難となる。
　従って、２回目の焼鈍後に、マルテンサイト相およびフェライト相の体積分率を制御して1180 MPa以上の引張強度を確保しつつ、所望の残留オーステナイト相の体積分率を確保して優れた延性を得るには、２回目の焼鈍後の冷却停止温度は350～450℃の範囲とする必要がある。 Cooling stop temperature: 350-450 ° C
Since the low temperature transformation phase from austenite becomes harder as the transformation temperature is lower, the steel plate residence temperature after cooling is stopped is important for controlling the strength of the low temperature transformation phase.
When the cooling stop temperature after the second annealing is lower than 350 ° C, it is cooled to a low temperature, so the low-temperature transformation phase is mainly a hard martensite phase, and it is easy to secure TS: 1180 MPa class, but the ductility is descend. On the other hand, when the cooling stop temperature after the second annealing is higher than 450 ° C., although the retained austenite phase is generated, the formation of the ferrite phase does not proceed, so that it becomes difficult to obtain excellent ductility.
Therefore, after the second annealing, the volume fraction of the martensite phase and the ferrite phase is controlled to ensure a tensile strength of 1180 MPa or more, while ensuring the desired volume fraction of the retained austenite phase and excellent ductility. In order to obtain this, the cooling stop temperature after the second annealing must be in the range of 350 to 450 ° C.
350～450℃での滞留時間：100～1000秒
　２回目の焼鈍後の冷却停止温度：350～450℃における滞留時間が100秒に満たないと、所望の残留オーステナイト体積分率を得ることが困難となり、滞留後の室温までの冷却過程において未変態のオーステナイト相がマルテンサイト相となり、マルテンサイト相の体積分率が過剰となる結果、高強度化するものの、延性が低下する。一方、1000秒を超えて滞留させると、残留オーステナイト相の生成が進行するため、延性は向上するものの、1180 MPa以上の引張強度を得るのが困難になる。それ故、1180 MPa以上の引張強度を確保すると共に、優れた延性を得るには、350～450℃における滞留時間は100～1000秒の範囲にする必要である。好ましくは200～800秒の範囲である。 Residence time at 350 to 450 ° C: 100 to 1000 seconds Cooling stop temperature after the second annealing: If the residence time at 350 to 450 ° C is less than 100 seconds, it is difficult to obtain the desired retained austenite volume fraction In the cooling process to room temperature after residence, the untransformed austenite phase becomes the martensite phase, and the volume fraction of the martensite phase becomes excessive. As a result, the strength is increased, but the ductility is lowered. On the other hand, when retained for more than 1000 seconds, the formation of a retained austenite phase proceeds, so that ductility is improved, but it becomes difficult to obtain a tensile strength of 1180 MPa or more. Therefore, in order to secure a tensile strength of 1180 MPa or more and to obtain excellent ductility, the residence time at 350 to 450 ° C. needs to be in the range of 100 to 1000 seconds. The range is preferably 200 to 800 seconds.
　冷却停止後の鋼板を上記滞留温度域に保持する手段としては、例えば、焼鈍後の冷却設備の下流工程に保温装置等を設けて、鋼板の温度を上記滞留温度に調整する手段等が挙げられる。なお、滞留後の鋼板は、従来公知の任意の方法により所望の温度に冷却される。 Examples of the means for holding the steel sheet after stopping the cooling in the residence temperature range include a means for adjusting the temperature of the steel sheet to the residence temperature by providing a heat retaining device or the like in the downstream process of the cooling equipment after annealing. . In addition, the steel plate after a residence is cooled to desired temperature by the conventionally well-known arbitrary methods.
　上記の焼鈍後、最終的に得られた冷延鋼板に、形状矯正や表面粗度調整の目的から調質圧延（スキンパス圧延）を行ってもかまわないが、過度にスキンパス圧延をすると鋼板に歪が導入されるため、結晶粒が展伸されて圧延加工組織となり、延性が低下するおそれがある。そのため、スキンパス圧延の圧下率は0.05％以上0.5％以下程度とすることが好ましい。 After the annealing described above, the cold-rolled steel sheet finally obtained may be subjected to temper rolling (skin pass rolling) for the purpose of shape correction or surface roughness adjustment. Therefore, the crystal grains are expanded to form a rolled structure, and the ductility may be reduced. Therefore, the rolling reduction of skin pass rolling is preferably about 0.05% to 0.5%.
　表１に示す成分組成になる鋼を溶製してスラブとし、1250℃に加熱後、仕上げ圧延機出側温度：900℃で熱間仕上げ圧延を施し、圧延終了後、80℃／秒の速度で冷却して、600℃で巻取り、ついで塩酸酸洗後、圧下率：40％の冷間圧延を施したのち、表２に示す条件で焼鈍処理および制御冷却処理を行い、板厚：1.6mmの冷延鋼板を製造した。なお、１回目の焼鈍時の冷却停止温度までの冷却速度は10～50℃／秒の範囲内の速度とした。
　得られた冷延鋼板について、下記に示す材料試験により材料特性を調査した。得られた結果を表３に示す。 Steel with the composition shown in Table 1 is melted into a slab, heated to 1250 ° C, hot finish rolling at the finishing mill exit temperature: 900 ° C, and after rolling is completed, the speed is 80 ° C / sec. After cooling at 600 ° C. and then pickling with hydrochloric acid, cold rolling with a reduction ratio of 40% was performed, and then annealing treatment and controlled cooling treatment were performed under the conditions shown in Table 2, and the plate thickness: 1.6 mm cold-rolled steel sheets were produced. The cooling rate up to the cooling stop temperature during the first annealing was set within a range of 10 to 50 ° C./second.
About the obtained cold-rolled steel plate, the material characteristic was investigated by the material test shown below. The obtained results are shown in Table 3.
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000001
Figure JPOXMLDOC01-appb-T000002
Figure JPOXMLDOC01-appb-T000002
（１）鋼板の組織
　圧延方向断面で、板厚の1/4位置の面を光学顕微鏡または走査型電子顕微鏡（SEM）で観察することにより調査した。観察はＮ＝５（観察視野５箇所）で実施した。フェライト相の体積分率は、倍率1000倍の断面組織写真および1000倍、3000倍の断面SEM写真を用い、画像解析により、任意に設定した10μm×10μm四方の正方形領域内に存在する相の占有面積を求め、これをフェライト相の体積分率とした。
　残留オーステナイトの量は、MoのKα線を用いてＸ線回折法により求めた。すなわち、鋼板の板厚1/4付近の面を測定面とする試験片を使用し、オーステナイト相の（２１１）面および（２２０）面と、フェライト相の（２００）面および（２２０）面とのピーク強度から残留オーステナイト相の体積率を算出した。
　各相の体積分率は、最初にフェライト相と低温変態相とに区別し、まずフェライト相の体積分率を決定し、次にＸ線により残留オーステナイト相の体積分率を決定し、残る体積分率をマルテンサイト相と判断した。
　また、マルテンサイト相の個数は、3000倍の断面SEM写真を用い、任意に設定した10μm×10μm四方の正方形領域内に存在するマルテンサイト相の数を求め、これをマルテンサイト相の数とした。 (1) Structure of steel sheet The cross section in the rolling direction was examined by observing the surface at 1/4 position of the sheet thickness with an optical microscope or a scanning electron microscope (SEM). The observation was carried out at N = 5 (5 observation fields). The volume fraction of the ferrite phase is the occupation of the phase existing in a square area of 10 μm × 10 μm square arbitrarily set by image analysis using a cross-sectional structure photograph of 1000 times magnification and cross-sectional SEM photographs of 1000 times and 3000 times The area was determined and used as the volume fraction of the ferrite phase.
The amount of retained austenite was determined by X-ray diffraction using Mo Kα rays. That is, using a test piece having a surface near a thickness of 1/4 of the steel sheet as a measurement surface, the (211) surface and (220) surface of the austenite phase, and the (200) surface and (220) surface of the ferrite phase From the peak intensity, the volume fraction of the retained austenite phase was calculated.
The volume fraction of each phase is first distinguished from the ferrite phase and the low temperature transformation phase, first the volume fraction of the ferrite phase is determined, then the volume fraction of the residual austenite phase is determined by X-rays, and the remaining volume The fraction was judged as the martensite phase.
In addition, the number of martensite phases was determined by using a cross-sectional SEM photograph of 3000 times the number of martensite phases present in an arbitrarily set 10 μm × 10 μm square area, and this was taken as the number of martensite phases. .
（２）引張特性
　圧延方向と90°の方向を長手方向（引張方向）とするJIS Z 2201に記載の５号試験片を用い、JIS Z 2241に準拠した引張試験を行い評価した。なお、引張特性の評価基準はTS×El≧22000MPa・％以上（TS：引張強度（MPa）、El：全伸び（％））を良好とした。 (2) Tensile properties A No. 5 test piece described in JIS Z 2201 with the rolling direction and 90 ° direction as the longitudinal direction (tensile direction) was used, and a tensile test based on JIS Z 2241 was performed and evaluated. The evaluation criteria for tensile properties were TS × El ≧ 22000 MPa ·% or more (TS: tensile strength (MPa), El: total elongation (%)).
Figure JPOXMLDOC01-appb-T000003
Figure JPOXMLDOC01-appb-T000003
　表３から明らかなように、発明例はいすれも、TS≧1180MPaで、かつTS×El≧22000MPa・％を満足する延性に極めて優れた高強度冷延鋼板が得られていることが分かる。
　これに対し、Ｃが本発明範囲を超えて多量に含有されたNo.６は所望の組織を発現することができず延性に劣っていた。
　また、製造条件が本発明範囲外である、１回目の焼鈍温度が低いNo.７、２回目の焼鈍温度が低いNo.11、２回目の焼鈍後の冷却速度が遅いNo.13および滞留時間が長いNo.18はいずれも、フェライト相の体積分率が高く、引張強度が1180MPa以上を満足していない。
　さらに、１回目の焼鈍温度が高いNo.８、１回目の焼鈍後の冷却停止温度が低いNo.９、１回目の焼鈍後の冷却停止温度が高いNo.10、２回目の焼鈍温度が高いNo.12、２回目の焼鈍後の冷却速度が速いNo.14、２回目の焼鈍後の冷却停止温度が低いNo.15および冷却停止温度が高いNo.16はいずれも、フェライト相の体積分率が少なく、延性に劣っていた。滞留時間が短いNo.17は、残留オーステナイト相の体積分率が少なく、延性に劣っていた。 As is apparent from Table 3, it can be seen that all the inventive examples have obtained high-strength cold-rolled steel sheets having excellent ductility satisfying TS ≧ 1180 MPa and satisfying TS × El ≧ 22000 MPa ·%.
On the other hand, No. 6 containing a large amount of C exceeding the range of the present invention could not express a desired tissue and was inferior in ductility.
In addition, the manufacturing conditions are outside the scope of the present invention, the first annealing temperature is low No. 7, the second annealing temperature is low No. 11, the cooling rate after the second annealing is slow No. 13 and the residence time No. 18, which has a long length, has a high volume fraction of ferrite phase and does not satisfy the tensile strength of 1180 MPa or more.
Furthermore, No. 8 with a high first annealing temperature, No. 9 with a low cooling stop temperature after the first annealing, No. 10 with a high cooling stop temperature after the first annealing, and a second annealing temperature high. No. 12, fast cooling rate after second annealing No. 14, low cooling stop temperature No. 15 after second annealing and No. 16 high cooling stop temperature are both the volume fraction of the ferrite phase The rate was low and the ductility was poor. No. 17 having a short residence time had a low volume fraction of retained austenite phase and was inferior in ductility.
　本発明に従い、鋼板中のＣ量を低減し、Nb，Cu、Ni、Cr、Mo、Ｖなど高価な元素を積極的に含有させずとも、フェライト相、マルテンサイト相および残留オーステナイト相各々の体積分率を規定することにより、安価でかつ優れた延性を有し、しかも引張強度（TS）が1180MPa以上の高強度冷延鋼板を得ることができる。また、本発明の高強度冷延鋼板は、自動車部品以外にも、建築および家電分野など厳しい寸法精度、加工性が必要とされる用途にも好適である。 According to the present invention, the volume of each of the ferrite phase, martensite phase and residual austenite phase is reduced without reducing the amount of C in the steel sheet and actively including expensive elements such as Nb, Cu, Ni, Cr, Mo and V. By specifying the fraction, it is possible to obtain a high-strength cold-rolled steel sheet that is inexpensive and has excellent ductility and has a tensile strength (TS) of 1180 MPa or more. The high-strength cold-rolled steel sheet of the present invention is also suitable for applications that require strict dimensional accuracy and workability, such as in the field of architecture and home appliances, in addition to automobile parts.

Claims (2)

  1. 　質量％で、
    　　Ｃ：0.16～0.20％、
    　　Si：1.0～2.0％、
    　　Mn：2.5～3.5％、
    　　Ｐ：0.030％以下、
    　　Ｓ：0.0050％以下、
    　　Al：0.005～0.1％、
    　　Ｎ：0.01％以下、
    　　Ti：0.001～0.050％および
    　　Ｂ：0.0001～0.0050％
    を含有し、残部がFe及び不可避的不純物からなる成分組成を有し、体積分率で、
    　　フェライト相：40～65％、
    　　マルテンサイト相：30～55％および
    　　残留オーステナイト相：５～15％
    を含み、圧延方向断面において単位面積：１μm2当たりのマルテンサイト相の数が 0.5～5.0個を満足する組織を有し、引張強度：1180MPa以上であることを特徴とする、延性に優れる高強度冷延鋼板。     % By mass
    C: 0.16-0.20%
    Si: 1.0-2.0%
    Mn: 2.5-3.5%
    P: 0.030% or less,
    S: 0.0050% or less,
    Al: 0.005-0.1%,
    N: 0.01% or less,
    Ti: 0.001 to 0.050% and B: 0.0001 to 0.0050%
    And the balance has a component composition consisting of Fe and inevitable impurities, with a volume fraction,
    Ferrite phase: 40-65%,
    Martensite phase: 30-55% and residual austenite phase: 5-15%
    With a structure satisfying 0.5 to 5.0 martensite phases per unit area: 1 μm 2 in the cross section in the rolling direction, and having a tensile strength of 1180 MPa or more, high strength with excellent ductility Cold rolled steel sheet.
  2. 　請求項１に記載の成分組成からなる鋼スラブを、熱間圧延し、ついで酸洗後、冷間圧延したのち、焼鈍を施して高強度冷延鋼板を製造するに際し、
    　上記焼鈍工程において、800～950℃で焼鈍後、冷却停止温度：200～500℃まで冷却し、ついで750～850℃に再加熱後、平均冷却速度：５～50℃/秒の速度で、350～450℃の冷却停止温度域まで冷却し、この温度域に100～1000秒滞留させることを特徴とする、延性に優れる高強度冷延鋼板の製造方法。     When the steel slab having the component composition according to claim 1 is hot-rolled, then pickled, cold-rolled, and then subjected to annealing to produce a high-strength cold-rolled steel sheet,
    In the above annealing process, after annealing at 800 to 950 ° C., the cooling stop temperature is cooled to 200 to 500 ° C., then reheated to 750 to 850 ° C., and the average cooling rate is 5 to 50 ° C./sec. A method for producing a high-strength cold-rolled steel sheet with excellent ductility, characterized by cooling to a cooling stop temperature range of up to 450 ° C. and retaining in this temperature range for 100 to 1000 seconds.
PCT/JP2012/002807 2012-04-24 2012-04-24 High strength cold-rolled steel plate of excellent ductility and manufacturing method therefor WO2013160938A1 (en)

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Citations (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2008169475A (en) * 2006-12-11 2008-07-24 Kobe Steel Ltd High-strength steel sheet
JP2011047042A (en) * 2009-07-29 2011-03-10 Jfe Steel Corp Process for production of high-strength cold-rolled steel sheet having excellent chemical conversion processability

Patent Citations (2)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP2008169475A (en) * 2006-12-11 2008-07-24 Kobe Steel Ltd High-strength steel sheet
JP2011047042A (en) * 2009-07-29 2011-03-10 Jfe Steel Corp Process for production of high-strength cold-rolled steel sheet having excellent chemical conversion processability

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