WO2001090431A1 - Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same - Google Patents

Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same Download PDF

Info

Publication number
WO2001090431A1
WO2001090431A1 PCT/JP2001/001004 JP0101004W WO0190431A1 WO 2001090431 A1 WO2001090431 A1 WO 2001090431A1 JP 0101004 W JP0101004 W JP 0101004W WO 0190431 A1 WO0190431 A1 WO 0190431A1
Authority
WO
WIPO (PCT)
Prior art keywords
less
cold
steel sheet
rolling
temperature
Prior art date
Application number
PCT/JP2001/001004
Other languages
French (fr)
Japanese (ja)
Inventor
Chikara Kami
Akio Tosaka
Takuya Yamazaki
Original Assignee
Kawasaki Steel Corporation
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
Priority claimed from JP2000156274A external-priority patent/JP4524859B2/en
Priority claimed from JP2000335803A external-priority patent/JP4665302B2/en
Application filed by Kawasaki Steel Corporation filed Critical Kawasaki Steel Corporation
Priority to KR1020027001080A priority Critical patent/KR20020019124A/en
Priority to DE60121162T priority patent/DE60121162T2/en
Priority to CA002379698A priority patent/CA2379698C/en
Priority to EP01906128A priority patent/EP1291448B1/en
Publication of WO2001090431A1 publication Critical patent/WO2001090431A1/en
Priority to US10/654,775 priority patent/US7101445B2/en
Priority to US10/655,361 priority patent/US7067023B2/en

Links

Classifications

    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • 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
    • C21D8/0421—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 working steps
    • C21D8/0426—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0236—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0273—Final recrystallisation annealing
    • 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
    • C22C38/001—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
    • C22C38/004—Very low carbon steels, i.e. having a carbon content of less than 0,01%
    • 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
    • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
    • 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
    • C21D1/185—Hardening; Quenching with or without subsequent tempering from an intercritical temperature
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/005—Ferrite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—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
    • C21D8/0421—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 working steps
    • C21D8/0436—Cold rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—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

Definitions

  • the present invention is used for parts where structural strength, especially strength and / or rigidity at the time of deformation is required, such as construction members, mechanical structural parts, and structural parts of automobiles.
  • Cold rolled steel sheet, electro-zinc-plated steel sheet, hot-dip galvanized steel sheet, and alloyed hot-dip galvanized steel sheet which are suitable as a material steel sheet of a molded body subjected to ascent heat treatment and have excellent strain age hardening characteristics It concerns the manufacturing method.
  • the method of softening before press-forming is to make the press-forming 'easy, and after press-forming, it is hardened to increase the part strength.
  • Nb is added in accordance with the content of C, N, and A1 in steel, and Nb / (dissolved C + dissolved N) is specified in at%.
  • the method of adjusting the solid solution C and solid solution N in the copper plate by controlling the cooling rate after annealing is described in Japanese Patent Publication No. 61-45689. Improves bake hardenability A method is disclosed.
  • the above steel sheet is made of a material having excellent deep drawability, the strength of the material steel sheet is low, and is not always sufficient as a structural material.
  • Japanese Patent Application Laid-Open No. 5-25549 discloses a method of improving bake hardenability by adding W, Cr, Mo alone or in combination to steel.
  • Japanese Patent Application Laid-Open No. 10-310847 discloses an alloyed hot-dip galvanized steel sheet whose tensile strength increases by 60 MPa or more in a heat treatment temperature range of 200 to 450. I have.
  • This steel sheet contains, in mass%, C: 0.01-0.08%, Mn: 0.01-3.0%, and one or more of W, Cr, and Mo in total of 0%. 0.05 to 3.0%, and if necessary, Ti: 0.005 to 0.1%, Nb: 0.005 to 0.1%, V: 0.005 to 0.1% Or, it has a composition containing two or more kinds, and the microstructure of the steel is mainly composed of ferrite or ferrite.
  • this technology forms fine carbides in the steel sheet by heat treatment after forming, effectively increases the dislocation with respect to the strain applied during pressing, and increases the amount of strain.
  • Heat treatment must be performed in the temperature range of 370 ° C, and the required heat treatment temperature is higher than the general bake hardening treatment temperature.
  • reduction of vehicle weight in automobiles has become a very important issue.
  • the strength of the steel plate used is increased, that is, the steel used is It is effective to make the board thinner.
  • Automotive parts using thin high-strength steel sheets must fully exhibit the characteristics appropriate to their role. Properties vary by part, but include, for example, dent resistance, static strength against bending and torsion, fatigue resistance, and impact resistance. In other words, high-strength steel sheets applied to automotive parts need to have excellent properties after forming. Since these characteristics are related to the strength of the steel sheet after forming, it is necessary to set the lower limit of the strength of the high-strength steel sheet to be used in order to achieve thinning.
  • a cold-rolled steel sheet for outer panel panels is known to use ultra-low carbon steel as a raw material and control the amount of C finally remaining in a solid solution in an appropriate range. ing.
  • This type of steel sheet is kept soft during press forming, secures shape freezing and ductility, and yield stress using the age hardening phenomenon that occurs in the paint baking process at 170 X 20 minutes after press forming is performed after press forming. It seeks to secure dent resistance by gaining an increase.
  • C forms a solid solution in the steel during press forming and is soft.
  • solid solution C adheres to dislocations introduced during press forming in the paint baking process. The yield stress increases.
  • JP-A-60-52528 discloses that C: 0.02 to 0.15%, Mn: 0.8 to 3.5%, P: 0.02 to 0.15%, A1: 0. 550 steel containing 10% or less, N: 0.005 to 0.025%.
  • a method for producing a high-strength thin steel sheet having both good hot rolling at a temperature of C or lower and controlled cold heat treatment of annealing after cold rolling and good ductility and spot weldability are both disclosed.
  • the steel sheet manufactured by the technique described in JP-A-60-52528 has a mixed structure composed of a low-temperature transformation product phase mainly composed of ferrite and martensite, has excellent ductility, and has a positive effect.
  • the aim is to obtain high strength by using strain aging during baking of paint with added N.
  • Japanese Patent Publication No. 5-24979 discloses that C: 0.08 to 0.20%, Mn: 1.5 to 3.5%, the composition of which consists of the balance Fe and unavoidable impurities.
  • a bake-hardenable high-tensile cold-rolled thin steel sheet composed of uniform bainite with a ferrite content of 5% or less or bainite partially containing martensite.
  • the structure is mainly made of payite by rapidly cooling the temperature range of 400 to 200 ° C in the cooling process after continuous annealing and then gradually cooling it. The aim is to obtain a higher bake hardening amount than ever before.
  • Japanese Patent Publication No. 61-12008 discloses a method for producing a high-strength steel sheet having a high r-value. As shown, this production method is characterized by annealing in the coexistence region of ferrite and austenite after cold rolling using ultra-low C steel as a material, and the resulting steel sheet has a high r value and high paint baking. It is said to have hardenability (BH properties), but the amount of BH obtained is at most about 60 MPa. Also, although the yield point of this steel sheet also increases after aging, there is no increase in TS, and the applicable parts are limited. There was a problem that there is.
  • the amount of C is set to 0.02 to 0.13 mass%, N is added to a large amount of 0.0080 to 0.0250 mass%, and then the finish rolling temperature and the winding temperature are controlled to control a large amount.
  • N is added to a large amount of 0.0080 to 0.0250 mass%, and then the finish rolling temperature and the winding temperature are controlled to control a large amount.
  • Japanese Patent Application Laid-Open No. 10-183301 discloses that, among the steel components, C and N are particularly limited to C: 0.01 to 0.12 mass%, N: 0.0001 to 0.01 mass%, and By controlling the crystal grain size to 8111 or less, high bake hardenability and high room temperature aging resistance can be achieved, while ensuring a high mass of 80 MPa or more and suppressing the AI amount to 45 MPa or less. Hot rolled steel sheets are presented.
  • the present invention has been developed in view of the above-described circumstances, and a cold-rolled steel sheet having excellent strain aging hardening characteristics, in which tensile strength is increased by press forming-heat treatment while maintaining excellent deep drawability during press forming. And to propose alloyed hot-dip galvanized steel sheets together with their advantageous production methods.
  • an excellent deep drawability with a TSX r value ⁇ 750 MPa and an excellent strain age hardening characteristic (BH ⁇ 80 MPa and ATS ⁇ 40 MPa) are used for deep drawing. It is to provide cold-rolled steel sheets and hot-dip galvanized steel sheets (including alloyed ones) together with their advantageous production methods.
  • the present invention solves the above-mentioned problems of the prior art, and has a soft and high formability and a stable quality characteristic suitable for an automobile part requiring a high degree of formability. It is easy to mold into parts, has no shape defects such as spring pack, torsion, warpage, cracks, etc.
  • heat treatment after molding into automobile parts provides sufficient strength as automobile parts, and High tensile cold-rolled steel sheets having a high r-value of 1.2 or more and excellent strain aging hardening properties that can sufficiently contribute to lightweight daggers, and these steel sheets can be manufactured industrially at low cost without disturbing the shape.
  • the purpose is to provide a manufacturing method. Disclosure of the invention
  • the present inventors manufactured steel plates with various composition and manufacturing conditions, and performed many material evaluation experiments.
  • N has been used as a strengthening element in areas where high workability is required. It has been found that by advantageously utilizing the large strain age hardening phenomenon developed by the action of elements, it is possible to easily achieve both improvement in formability and high strength after molding. Furthermore, the present inventors have found that in order to advantageously utilize the strain age hardening phenomenon due to N, the strain age hardening phenomenon due to N is advantageously combined with the baking conditions of automobiles or the heat treatment conditions after molding more positively.
  • the present inventors further reduced the C content, performed continuous annealing at a temperature in the two-phase region of the fly-to-stenite, and controlled the subsequent cooling to reduce the content in the fly phase.
  • the microstructure containing 5% or more of the ferrite phase in the area ratio is high, and the combination of such a microstructure and an appropriate amount of dissolved N has a high r value, excellent press formability, and distortion. It has been found that a cold rolled steel sheet having excellent age hardening characteristics can be obtained. In addition, they have found that N can be fully utilized without the problem of room-temperature aging degradation, which was a conventional problem.
  • the present inventors use N as a strengthening element, control the A1 content in an appropriate range according to the N content, and optimize the hot rolling conditions, cold rolling, and cold rolling annealing conditions, By optimizing the visual structure and solid solution N, the r-value is much higher than the conventional solid solution strengthened C-Mn steel sheet and precipitation hardened It has been found that a steel sheet having strain aging hardening characteristics not found in the above steel sheets can be obtained.
  • the steel sheet of the present invention has a higher strength after a paint baking treatment by a simple tensile test than a conventional steel sheet, and has a small variation in strength when plastically deformed according to actual pressing conditions, and is a stable component.
  • the strength characteristics can be obtained, and it can be applied to parts that require reliability. For example, the part where the plate thickness is reduced due to large strain is larger than other parts, and the hardening allowance is more uniform when evaluated by the load capacity of (plate thickness) X (strength). Direction, and the strength as a part is stable.
  • solid solution N has a larger interaction with dislocations introduced by forming, even if the heat treatment temperature after forming is lowered, even if solid solution N reaches higher yield stress than solid solution C.
  • the dislocations introduced in the predeformation are difficult to move.
  • the present invention is based on the above findings.
  • the above findings were obtained from the following experiments. At mass%, C: 0.0015%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0.008%, and N: 0 0.11%, and Nb 0.005 to 0.05% and A1 0.0 Contained in the range of 05-0.03%, the balance is Fe and unavoidable impurities.
  • the composition of the sheet-pattern (thickness: 30 mm) is 115 (after uniform heating with TC, the finishing temperature is 900 with Ar 3 transformation point or higher). Hot rolling was performed in three passes so as to obtain C. After the rolling was completed, water cooling was performed 0.1 seconds later, and then heat treatment equivalent to coil winding was performed at 500 for 1 hour.
  • the obtained hot-rolled sheet having a thickness of 4 mra was cold-rolled at a rolling reduction of 82.5%, then recrystallized and annealed at 800 C for 40 seconds, and then temper-rolled at a rolling reduction of 0.8%.
  • a JIS No. 5 tensile test piece was sampled in the rolling direction, and the tensile strength was measured at a strain rate of 0.02 / s using an ordinary tensile tester.
  • JIS No. 5 tensile test specimens sampled in the rolling direction from these cold-rolled sheets were given a 10% tensile strain, heat-treated for 120 and 20 minutes, and then subjected to a normal tensile test. .
  • Figure 1 shows the results of a study on the relationship between the steel composition (N%-14/93 ⁇ Nb%-14/27 ⁇ Al%-14/11 ⁇ B%) and ATS.
  • the obtained hot-rolled sheet having a thickness of 4 ram was cold-rolled at a reduction ratio of 82.5%, and then 820.
  • C recrystallization annealing for 40 seconds, followed by temper rolling at a reduction of 0.8%.
  • Figure 2 shows the results of a study on the relationship between the B content in steel and ATS. As shown in the figure, when B is contained in 0.0005 to 0.0015 mass%, a high ⁇ TS of 60 MPa or more can be obtained.
  • the B content is less than 0.0005 mass%, the effect of refining crystal grains by adding Nb in combination is small. Conversely, if the B content exceeds 0.0015 mass%, the amount of B segregating in the vicinity of the grain boundary increases, and such B atoms have a strong interaction with N atoms, so that effective B It is probable that ATS decreased due to the decrease in N content.
  • the resulting hot-rolled sheet with a thickness of 4 nun is cold-rolled at a rolling reduction of 82.5%, then recrystallized and annealed at 8803 ⁇ 4 for 40 seconds, and then temper-rolled with a rolling reduction of 0.8%.
  • a rolling reduction of 82.5% is cold-rolled at a rolling reduction of 82.5%, then recrystallized and annealed at 8803 ⁇ 4 for 40 seconds, and then temper-rolled with a rolling reduction of 0.8%.
  • a JIS No. 5 tensile test specimen was sampled in the rolling direction, and the tensile strength was measured at a strain rate of 0.02 / s using an ordinary tensile tester. Separately, a 10% tensile strain was applied to tensile test specimens collected from these cold-rolled sheets, and the specimens were treated at various temperatures for 20 minutes. After the heat treatment, the steel sheet was subjected to a normal tensile test.
  • Figure 3 shows the results of an investigation on the effect of the post-molding heat treatment temperature on.
  • steel A which is a steel with extremely low carbon and high N content
  • steel B which is a semi-ultra low carbon and low N steel. It shows high ATS, and shows similar ATS at high temperature. From these experimental results, it is necessary to use solid solution N to secure ATS at low temperatures.
  • Fig. 4 shows the effect of crystal grain size d and steel composition (N%-14/93-Nb%-14 /) on the reduction in elongation due to normal temperature aging ( ⁇ ⁇ 1) and the increase in tensile strength after forming ( ⁇ TS). 27-Al%-14/11. B%).
  • the amount of decrease in elongation ( ⁇ 1) was calculated using the total elongation measured with a JIS No. 5 test piece taken in the rolling direction from the cold-rolled sheet and 100, which is a normal temperature aging acceleration treatment using a separately taken test piece. The evaluation was made based on the difference from the total elongation measured after the time holding treatment.
  • the value of (N%-14/93-Nb%-14/27-Al%-14 / 11B%) is 0.0015 mass% or more and the crystal grain size d is 20 / m or less. It can be seen that in the case of, both high ATS and low ⁇ 1 can be achieved.
  • a No. 5 tensile test piece was sampled and subjected to a tensile test using an ordinary tensile tester at a strain rate of 3 ⁇ 10— / s, and the TSX r value, BH, and ATS were measured.
  • Figure 6 shows the relationship between these measured values and the B content.
  • B In addition to BH ⁇ 80MPa in the range of 0.0003% to 0.0015%, ⁇ TS levels of ⁇ TS ⁇ 60MPa and TSX r value ⁇ 850MPa, which are higher than those of B and 0.0003%, were achieved.
  • a cold-rolled steel sheet having excellent strain aging hardening characteristics characterized in that the balance is composed of Fe and unavoidable impurities.
  • N 0.005 to 0.440%
  • N / Al 0.30 or more
  • a cold-rolled steel sheet having excellent strain aging hardening characteristics characterized in that the balance is composed of Fe and unavoidable impurities.
  • one or more of Cu, Ni, and Mo be contained in a mass% of 1.0% or less as necessary.
  • the steel sheet preferably has a crystal grain size of 20 / zm or less.
  • the strength increase after molding is 60 MPa or more in the low temperature range of heat treatment temperature: 120 to 200.
  • the surface of the cold-rolled steel sheet may be provided with an electrogalvanized layer, a hot-dip galvanized layer, and an alloyed hot-dip galvanized layer.
  • the second invention is based on mass%
  • A1 0.005 to 0.030%
  • N 0.005 to 0.404%
  • the steel slab having a composition of substantially Fe is hot-rolled. At that time, cooling is started immediately after finishing rolling, and the winding temperature is 400 to 800. C rolled, then cold-rolled at a reduction rate of 60-95%, and then re-crystallized at 650-900 to produce a cold-rolled steel sheet with excellent strain age hardening characteristics. Manufacturing method.
  • the temperature of the temperature range up to the recrystallization temperature to 500 at a rate of 1 ⁇ 20 e C / s.
  • a hot-dip galvanizing treatment and then a heat alloying treatment may be performed.
  • the third aspect of the present invention provides a
  • n 0.01 to; 1.5%
  • A1 0.005 to 0.002%
  • N 0.0050 to 0.040%
  • TS Xr value Cold-rolled steel sheet for deep drawing with excellent strain aging hardening characteristics characterized by being 750 MPa or more.
  • V 0.005 to 0.10%
  • Solid solution N 0.0010% or more
  • the fourth invention is based on mass%
  • n 0.01-1.0%
  • P 0.1% or less
  • V 0.005 to 0.10%
  • N / A1 and 0.0010% or more of N as a solid solution, with the balance being Fe and unavoidable impurities.
  • Formability, strain aging hardening characteristics and room temperature aging resistance characterized by having an organization consisting of a ferrite phase and a ferrite phase having an average crystal grain size of 20 / xm or less and an r value of 1.2 or more. It is a high-tensile cold-rolled steel sheet with excellent resistance
  • one or more of the following groups a to c be further contained in mass% in addition to the composition.
  • Group a 1.0% or less in total of one or more of Cu, Ni, Cr, and Mo Group b: 0.1% or less in total of one or two of Ti and V
  • Group c One or two of Ca and RE are 0.0010 to 0.010% in total.
  • n 2.0% or less
  • a steel slab containing one or more of the following and having a N / A1 of 0.3 or more is heated to a slab heating temperature of 1000 ° C or more, Rough rolling and sheet par, none
  • Winding temperature hot rolling process at 650 ° C or lower to make hot rolled sheet
  • the cooling rate is 10 to 300 / s to a temperature range of 500 ° C or less.
  • r-value This is a method for producing cold-rolled steel sheets with excellent formability, strain aging hardening properties, and normal-temperature aging resistance, having at least 1.2.
  • one or more of the following groups a to c be further contained in mass% in addition to the composition.
  • Group a 1.0% or less in total of one or more of Cu, Ni, Cr, Mo, Group b: 0.1% in total of one or more of Ti, V Less than
  • Group c Ca, REM 1 or 2 types in total 0.0010 ⁇ 0.010%
  • the 7th Honoki is raass%
  • N / A1 and 0.0010% of N as a solid solution, with the balance being Fe and unavoidable impurities, and a ferrite with an average grain size of 10 m or less.
  • Group d One or more of Cu, Ni, Cr, and Mo is 1.0% or less in total.
  • Group e One or more of Nb, Ti, and V is 0.1 in total. %Less than
  • f group: B is less than 0.0030%
  • g group Ca or REM 1 or 2 kinds in total 0.0010 to 0.010%
  • the eighth present invention provides, in mass%,
  • Annealing temperature is applied to the cold-rolled sheet at a temperature not lower than the recrystallization temperature and not higher than 800 ° C.
  • a continuous annealing process is performed at the transformation point of Ac1 to (transformation point of Ac3-20 ° C), and thereafter, the cold-rolled sheet annealing process of cooling at a rate of 10 to 300 / s to a temperature range of 500 ° C or less is performed sequentially.
  • an overaging treatment for a residence time of 20 s or more is performed in a temperature range of 350 ° C. or less to the cooling stop temperature of the cooling.
  • one or more of the following groups d to g be further contained in mass% in addition to the composition.
  • Group d One or more of Cu, Ni, Cr and Mo are 1.0% or less in total,
  • f group: B is less than 0.0030%
  • g group Ca or REM 1 or 2 kinds in total 0.0010 to 0.010%
  • Figure 1 is a graph showing the relationship between the steel composition (N%-14/93-Nb%-14/27-Al%-14/11-B%) and the post-forming tensile strength rise (ATS).
  • Fig. 2 is a graph showing the relationship between B content and ATS in Nb and B composite added steel.
  • Fig. 3 is a graph comparing the difference in tensile strength rise between steel B with a large amount of solid solution C (conventional steel) and steel A with a large amount of solid solution N (inventive steel) due to post-forming heat treatment at a low temperature range. is there.
  • Figure 4 Grain size d and steel composition (N%-14/93 ⁇ Nb%-14/27 ⁇ Al) affecting the amount of decrease in elongation due to normal temperature aging ( ⁇ 1) and increase in tensile strength after forming (ATS) %-14/11 ⁇ B%).
  • Figure 5 is a graph showing the relationship between TSX i fi, BH, ATS and N / (A1 + Nb + B).
  • FIG. 6 is a graph showing the relationship between the TSX r value, ⁇ , ⁇ TS, and the B amount.
  • the amount of C is preferably suppressed to less than 0.01 mass%. It is more preferably at most 0.0050 mass%, more preferably at most 0.0030 mass%.
  • Si is a useful component that suppresses the decrease in elongation and improves the strength.However, if the content is less than 0.005 mass%, the effect of its addition is poor, and if it exceeds 1.0 mass%, the surface properties will be deteriorated. Therefore, Si was limited to the range of 0.005 to 1.0 mass% because it deteriorated the ductility. More preferably, it is in the range of 0.01 to 0.75 mass%.
  • Mn is not only useful as a strengthening component of steel, but also has the effect of forming Mn S to suppress embrittlement due to S. However, if the content is less than 0.01 mass%, the effect of its addition is poor. On the other hand, if the content exceeds 1.5 mass%, the surface properties deteriorate and the ductility decreases. Therefore, Mn is contained in the range of 0.01 to 1.5 mass%. More preferably, it is 0.10 to 0.75 mass%. P: 0.10 mass% or less
  • P effectively contributes to the strengthening of steel as a solid solution strengthening component, but when added in excess of 0.10 mass%, deep drawability is reduced due to the formation of phosphides such as (FeNb) xP. Therefore, P was limited to 0.10 mass% or less.
  • Al is added as a deoxidizing agent and to improve the yield of carbonitride-forming components.However, if the content is less than 0.005 mass%, it has no sufficient effect, whereas if it exceeds 0.003 mass%, steel is added. This leads to an increase in the amount of N to be added into the steel, which tends to cause slab defects during steelmaking. Therefore, A1 is contained in the range of 0.005 to 0.30 mass%.
  • N is an important element that plays a role in imparting strain age hardening characteristics to a steel sheet in the present invention.
  • the content is less than 0.005 mass%, sufficient strain aging hardening properties cannot be obtained, while addition of a large amount exceeding 0.004 mass% causes a decrease in press formability. Therefore, N was contained in the range of 0.005 to 0.440 mass%. It is more preferably 0.008 to 0.015 mass%.
  • B when added in combination with Nb, has the effect of effectively refining the hot-rolled structure and the cold-rolled recrystallized structure and improving the secondary work brittleness resistance.
  • the content is less than 0.0001 mass%, a sufficient refining effect cannot be obtained.
  • the content exceeds 0.003 mass%, not only does the amount of BN precipitate increase, but also hinders solution formation in the slab heating step. It will be. Therefore, B is contained in the range of 0.0001 to 0.003 mass%. It is more preferably 0.0001 to 0.0015 mass%, and more preferably 0.0007 to 0.0012raass%.
  • Nb 0.005 to 050 mass%
  • Nb contributes effectively to the refinement of the hot-rolled structure and the cold-rolled recrystallization-annealed structure by the combined addition with B, and has the effect of fixing solid solution C as NbC.
  • Nb forms a nitride called NbN, which contributes to the refinement of the cold-rolled recrystallization annealing structure.
  • Nb content is less than 0.005 raass%, it becomes difficult not only to precipitate and fix solid solution C, but also to make the hot-rolled structure and the cold-rolled recrystallization annealed structure insufficiently refined.
  • it exceeds 0.050 mass% ductility is reduced.
  • Nb is contained in the range of 0.005 to 0.050 mass%. Preferably, it is 0.010 to 0.030 raass%. Further, as described above, Nb has an effect of fixing solid solution C as NbC. Also, a nitride such as NbN is formed. Similarly, A1 and B form A1N and BN, respectively. Therefore, it is important to satisfy the relations of the following equations (1) and (2) in order to sufficiently secure the dissolved N amount and sufficiently reduce the solute C.
  • Steel having the above-mentioned preferred composition is smelted by a known smelting method such as a converter, and is made into a slab by an ingot-making method or a continuous sintering method.
  • the heating temperature of hot rolling is not particularly specified, but it is advantageous to fix solid solution C and precipitate as carbide in order to improve the deep drawability.
  • the hot rolling temperature is preferably set to 130 (TC or lower. In order to further improve the workability, the temperature is preferably set to 1150 ° C. or lower. However, the heating temperature is 900 ° C. If the temperature is less than C, the improvement in workability is saturated, and conversely, the rolling load during hot rolling increases and the risk of rolling trouble increases, so the lower limit of the heating temperature is preferably 900.
  • the total reduction in hot rolling is preferably set to 70% or more. This is because if the total draft is less than 70%, the grain refinement of the hot-rolled sheet will be insufficient.
  • the finish rolling in the hot rolling is preferably completed in a temperature range of 960 to 650 ° C, and the hot rolling finish temperature may be in the y range above the Ar 3 transformation point but in the ⁇ range below the Ars transformation point. It may be. If the hot rolling finish temperature exceeds 0%, the crystal grains of the hot rolled sheet become coarse, and the deep drawability after cold rolling and annealing deteriorates. On the other hand, if it is less than 650, the deformation resistance increases, so that the hot rolling load increases and rolling becomes difficult.
  • the above cooling treatment conditions are not particularly limited, but the cooling start time is preferably within 1.5 seconds, more preferably within 1.0 seconds, and more preferably after the finish rolling. It is desirable to keep it within 0.5 seconds. This is because cooling immediately after the end of rolling increases the degree of supercooling with accumulated strain, so that more ferrite nuclei are generated, ferrite transformation is accelerated, and solidification in the ⁇ phase is promoted. This is because diffusion of the melt into the ferrite grains is suppressed, and the amount of solid solution present at the ferrite grain boundaries increases.
  • the cooling rate is preferably set to 10 ° C / s or more in order to secure solid solution.
  • the hot-rolling finishing temperature is equal to or higher than the Ar 3 transformation point, it is more preferable to set the cooling rate to 50 ° C / s or higher in order to secure solid solution N.
  • the hot rolled sheet is wound up into a coil.
  • it exceeds 800 ° C the scale formed on the surface of the hot-rolled sheet becomes thicker, which not only increases the load of scale removal work, but also increases the load. nitridation proceeds leads to variation in the coil longitudinal direction of solute N amount, while in ⁇ Ri temperature is less than 400 e C, since the flame Certificates up work piece, ⁇ Ri temperature of the hot-rolled sheet is 800 It must be in the range of 400 ° C.
  • the hot rolled sheet is subjected to cold rolling, and the rolling reduction in such cold rolling needs to be 60 to 95%. This is because a high r value is expected when the rolling reduction of cold rolling is less than 60%. No, on the other hand, if it exceeds 95%, the r-value will decrease.
  • the cold rolled sheet that has been subjected to the above cold rolling is then subjected to recrystallization annealing.
  • the annealing method may be either continuous annealing or patch annealing, but continuous annealing is more advantageous.
  • the continuous annealing may be either a normal continuous annealing line or a continuous hot-dip galvanizing line.
  • the annealing conditions are preferably 650 t or more and 5 seconds or more. This is because if the annealing temperature is less than 650 C and the annealing condition is less than 5 seconds, recrystallization is not completed, and the deep drawability decreases. In order to further improve the deep drawability, it is desirable to perform annealing in a ferrite single phase region of 800 or more for 5 seconds or more.
  • annealing in the ⁇ + y two-phase region at a higher temperature partially causes the o: ⁇ y transformation to develop ⁇ 1 1 1 ⁇ texture and improve the r value, but the a ⁇ y transformation is completely If it progresses, the texture will be randomized, so the r-value will decrease and the deep drawability will be impaired.
  • the upper limit of the annealing temperature is 900. This is because, when the annealing temperature exceeds 900, the re-dissolution of the carbide proceeds and the solid solution C excessively increases, so that the delayed aging property decreases, and when the ⁇ - ⁇ transformation occurs This is because the texture is randomized, so that the r-value decreases and the deep drawability is impaired.
  • the temperature range from 500 ° C to the recrystallization temperature is gradually heated, and A1N etc. are sufficiently precipitated to effectively reduce the crystal grain size of the steel sheet. can do.
  • the temperature range in which the above-described controlled heating should be performed is from 500 ° C. at which A1N or the like starts to precipitate to the recrystallization temperature.
  • the heating rate is preferably in the range of 1 to 20 ° C / s. This is because if the heating rate is higher than 20 ° C, a sufficient amount of precipitate cannot be obtained, while if it is lower than l ° C / s, the precipitate becomes coarse and the effect of suppressing the grain growth is weakened.
  • the cooling rate after soaking in the recrystallization annealing is preferably set to 10 to 5 ° C / s.
  • a cooling rate of 10 ° CZs or less grain growth occurs during cooling, crystal grains become coarse, and strain aging characteristics and aging characteristics at room temperature decrease.
  • a cooling rate of 10 ° CZs or less grain growth occurs during cooling, crystal grains become coarse, and strain aging characteristics and aging characteristics at room temperature decrease.
  • diffusion of N in the solid solution state to the grain boundary does not occur sufficiently, and the aging characteristics at room temperature deteriorate. Still more preferably, it is 10 to 3 OtZs.
  • a hot-dip galvanizing treatment followed by a heat alloying treatment is performed to form an alloyed hot-dip galvanized steel sheet.
  • the hot-dip galvanizing treatment and alloying treatment there is no particular limitation on the hot-dip galvanizing treatment and alloying treatment, and the treatment may be performed according to a conventionally known method.
  • steel sheets with alloyed hot-dip galvanized steel and then temper-rolled to improve workability and appearance after processing (dull-finished steel sheet, bright-finished steel sheet, and a specific roughness pattern on the surface)
  • Steel sheet which has been subjected to surface treatment that is usually used as a thin steel sheet, such as a steel sheet having an oil layer such as a fireproof oil or a lubricating oil on the surface. Effect can be fully enjoyed.
  • the amount of C is preferably suppressed to less than 0.01 mass%. More preferably, it is not more than 0.0050raass%, more preferably not more than 0.0030raass%. From the viewpoint of securing strength and preventing coarsening of crystal grains, C is 0.0. It is desirable to contain 005% or more.
  • Si is a useful component that suppresses the decrease in elongation and improves the strength.However, if the content is less than 0.005 mass%, the effect of its addition is poor, while if it exceeds 1.0 mass%, the surface properties will deteriorate. Therefore, Si was limited to the range of 0.005 to 1.0 mass% because it deteriorated the ductility. More preferably, it is in the range of 0.01 to 75 mass%.
  • Mn is not only useful as a strengthening component of steel, but also has the effect of forming Mn S to suppress embrittlement due to S. However, if the content is less than 0.01 mass%, the effect of its addition is poor. On the other hand, if the content exceeds 1.5 mass%, deterioration of the surface properties and decrease in ductility are caused. Therefore, Mn is contained in the range of 0.01 to 1.5 mass%. More preferably, it is 0.10 to 0.75raass%.
  • P effectively contributes to the strengthening of steel as a solid solution strengthening component.
  • phosphides such as (FeNb) xP are formed, and the deep drawability decreases. Therefore, P was limited to 0.10 mass% or less.
  • A1 is added as a deoxidizing agent and to improve the yield of carbonitride-forming components.However, if the content is less than 0.005% by mass, it has no sufficient effect, whereas if it exceeds 0.003% by mass, the addition of A1 This leads to an increase in the amount of N to be added into the steel, which tends to cause slab defects during steelmaking. Therefore, A1 is contained in the range of 0.005 to 0.30 mass%.
  • N is an important element that plays a role in imparting strain age hardening characteristics to a steel sheet in the present invention.
  • the content is less than 0.005 mass%, sufficient strain aging
  • the addition of a large amount exceeding 0.040 mass% leads to a decrease in press formability. Therefore, N was contained in the range of 0.005 to 0.040 mass%. Preferably, it is 0.008 to 0.015 mass%.
  • B when added in combination with Nb, has the effect of effectively refining the hot-rolled structure and the cold-rolled recrystallized structure and improving the secondary work brittleness resistance.
  • the content is less than 0.0001 mass%, sufficient refining effect cannot be obtained.
  • the content exceeds 0.003 mass%, not only the BN precipitation amount increases, but also the solution solution at the slab heating stage may be hindered. become. Therefore, B was contained in the range of 0.0001 to 0.003 raass%. It is more preferably 0.0001 to 0.0015 mass%, more preferably 0.0007 to 0.0012 mass%.
  • Nb 0.005 to 0.050%
  • Ti 0.005 to 0.070%
  • V 0.005 to 0.10%
  • Nb, Ti, and V contribute to the refinement of the hot-rolled structure and the cold-rolled recrystallized structure when combined with B, and have the function of precipitating solid solution C as NbC, TiC, and VC. If necessary, they are added together with B, but if each is less than 0.005%, their function is insufficient. On the other hand, if Nb exceeds 0.050%, Ti exceeds 0.070%, and V exceeds 0.10%, ductility deteriorates. Therefore, Nb was 0.005 to 0.050%, Ti was 0.005 to 0.070%, and V was 0.005 to 0.10%. Further, as described above, Nb has an effect of fixing solid solution C as NbC. Also, a nitride such as NbN is formed. Similarly, A1 and B form A1N and BN, respectively. Therefore, it is important to satisfy the following formulas (1) and (2) in order to secure a sufficient amount of solute N and sufficiently reduce the amount of solute C.
  • A1 forms A1N to reduce solute N.
  • N / A1 must be 0.30 or more.
  • NbN, TiN, VN and BN are formed to reduce the amount of solute N, so that the appropriate amount of solute N should be secured.
  • (Al + Nb + Ti + V + B) must be 0.30 or more.
  • Solid solution N 0.0010% or more
  • solid solution N is present at a content of 0.0010% or more.
  • the amount of solute N is determined by subtracting the amount of precipitated N from the total amount of N in the steel.
  • the electrolytic extraction analysis method using the constant potential electrolysis method.
  • the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides.
  • Electrolysis is performed at a constant potential using an acetyl-aceton system as an electrolyte.
  • the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed the best correspondence with the actual component strength.
  • the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N.
  • the amount of solid solution N is preferably 0.0015% or more, more preferably 0.0020% or more, and still more preferably 0.0030% or more.
  • the cold-rolled steel sheet of the present invention has the above composition and a TSX r value ⁇ This is a cold-rolled steel sheet for deep drawing, which is characterized by having excellent strain age hardening characteristics.
  • the TSXr value is less than 750MPa, it cannot be widely applied to structural members. Further, in order to further expand the applicable range, the TSXr value is preferably set to 850 MPa or more.
  • nO. ⁇ X SOiuin is adopted as a standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, hardening is achieved even with milder (lower temperature) treatment, in other words, aging conditions can be broadened. And it is possible. In general, in order to increase the amount of aging, it is advantageous to hold at a higher temperature and for a longer time as long as the material is not softened by excessive aging.
  • the lower limit of the heating temperature at which hardening is remarkable after pre-deformation is approximately 100.
  • the heating temperature exceeds 300 ° C, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous.
  • the holding time is about 30 s or more when the heating temperature is about 200 ° C, almost sufficient curing can be achieved.
  • the holding time is preferably 60 s or more. However, holding for more than 20 min not only does not allow further curing, but also significantly reduces production efficiency, which is disadvantageous in practical use.
  • the aging treatment conditions were evaluated at a heating temperature of 170 ° C. and a holding time of 20 min under the conventional paint baking treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention.
  • the method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used. Alternatively, only the portion where the strength is to be increased may be selectively heated.
  • the strength of automotive components needs to be able to withstand complex external stress loads, so that not only the strength characteristics in a small strain range but also the strength characteristics in a large strain range are important for a material steel plate.
  • the present inventors set the BH of the steel sheet of the present invention, which is to be used as a material for automobile parts, to be 80 MPa or more and the ATS to be 40 MPa or more. More preferably, the pressure is BHIOOMPa or more, and the temperature TS50MPa or more.
  • the heating temperature and / or the holding time during the aging treatment may be set to a higher temperature and / or a longer time.
  • the steel sheet of the present invention has an unprecedented advantage that it does not undergo aging deterioration (phenomenon in which YS increases and E 1 decreases) even when it is left at room temperature for a long period of about one year when it is not formed. Equipped. Further, in the present invention, there is no problem even if hot dip galvanizing or alloyed hot dip galvanizing is performed on the surface of the above-mentioned cold rolled steel sheet of the present invention, and TS, BH, and mu TS are comparable to those before plating. .
  • any of electroplating, electroplating, electrochromic plating, and electroplating can be preferably used. Fourth manufacturing conditions according to the present invention will be described.
  • C less than 0.01%
  • N 0.0050 to 0.04%
  • A1 0.005 to 0.03%
  • Si 0.005 to 1.0%
  • Mn 0.01 to 1.5%
  • P 0.1% or less
  • S 0.01% or less
  • B 0.001 to 0.003%
  • Nb 0.005 to 0.050%
  • Ti 0.005 to 0.070%
  • V 0.005 to 0.10%
  • This steel material is heated and soaked, then hot-rolled into a hot-rolled sheet. If the heating temperature (SRT) is too low, the effect of improving workability saturates, and the rolling load during hot rolling increases, which may cause rolling troubles and may cause insufficient uniformity of solid solution N.
  • the SRT should be 950 or higher.
  • the SRT is 1300 and the following is preferable.
  • 1150. C or less is preferable.
  • the grain refinement of the hot-rolled sheet becomes insufficient, so that it is preferably at least 80%.
  • finish rolling is completed in a temperature range exceeding ⁇ ⁇ 3, the texture becomes random due to ⁇ ⁇ ⁇ transformation, and excellent deep drawability cannot be obtained.
  • finish rolling is preferably performed in a temperature range of Ars to 600 or more.
  • finish rolling If lubrication rolling is not performed during finish rolling, additional shearing force acts on the surface of the steel sheet due to the frictional force between the roll and the steel sheet. ⁇ Since the orientation is preferentially formed, deep drawability deteriorates. Therefore, the finish rolling is preferably performed while lubricating.
  • the hot rolled sheet is wound into a coil.
  • the material to be processed after the winding step is also called a coil.
  • the winding temperature (CT) of the hot-rolled sheet is higher at higher temperatures, which is advantageous for coarsening of carbides.
  • CT winding temperature
  • it exceeds 800 the scale formed on the hot-rolled sheet surface becomes thicker and the load of scale removal work increases. Or the formation of nitrides progresses, causing a change in the amount of solute N in the longitudinal direction of the coil.
  • CT is preferably set to 800 to 400 t.
  • the obtained hot rolled sheet is subjected to recrystallization annealing by continuous annealing or patch annealing.
  • This annealing (hot rolled sheet annealing) is performed in order to obtain a recrystallized texture by recrystallizing the rolled texture formed by the ⁇ region warm rolling performed in the finish rolling.
  • the hot-rolled sheet is cold-rolled to be a cold-rolled sheet. If the rolling reduction of the cold rolling is less than 60%, a high r-value cannot be expected, while if it exceeds 95%, the r-value rather decreases, so that it is preferable to be 60 to 95%.
  • the cold rolled sheet is subjected to recrystallization annealing.
  • This annealing is preferably performed in either a continuous annealing line or a continuous molten zinc plating line.
  • the annealing conditions are preferably annealing temperature of 650 or more and X holding time of 5 seconds or more. Unless any of the annealing temperature of 650 or more and the holding time of 5 seconds or more are satisfied, recrystallization is not completed and the deep drawability decreases.
  • the annealing temperature is preferably 800 ⁇ or more and the X holding time is preferably 5 seconds or more.
  • the annealing temperature exceeds 900, the re-dissolution of carbides proceeds and the solute C increases excessively, so that the slow aging (normal temperature aging resistance) is reduced and the ⁇ ⁇ ⁇ transformation occurs. Since the texture is randomized and the r-value is reduced and the deep drawability is impaired, the annealing temperature is 900 It is preferable that the temperature be below ° C.
  • the cold-rolled annealed sheet obtained by recrystallizing and annealing the cold-rolled steel sheet is subjected to hot-dip galvanizing or further alloying as necessary. It is preferable that the cooling rate from the time before to the plating process is 5 / s or more, and the sheet temperature when hot-dip galvanizing is 400 to 600 ° C. In the alloying process, the processing temperature is 400 to 600 ° C. The processing time is preferably 5 to 40 seconds.
  • the cold-rolled steel sheet or the hot-dip galvanized steel sheet after the recrystallization annealing may be subjected to temper rolling for shape correction and surface roughness adjustment.
  • the rolling reduction of this temper rolling is preferably 10% or less. This is because when the rolling reduction exceeds 10%, the r value decreases. The reason for limiting the composition of the high-tensile cold-rolled steel sheet of the fifth invention will be described.
  • c must be contained at least 0.0015% in order to control the structure uniformly and finely and to secure a sufficient amount of the ferro-ferrite phase.
  • the content exceeds 0.025%, the carbide fraction in the steel sheet becomes excessive, and the ductility, the r value, and the formability are significantly reduced.
  • C is limited to the range of 0.0015 to 0.025%.
  • the content is preferably at most 0.020%, more preferably at most 0.010%.
  • the C content exceeds (12/93) Nb (%) (where Nb is the Nb content (%)).
  • Si is a useful element that can increase the strength of a steel sheet without significantly reducing the ductility of the steel, and is preferably contained in the present invention in an amount of 0.005% or more, particularly when high strength is required. Is more preferably 0.10% or more.
  • Si significantly raises the transformation point during hot rolling, making it difficult to secure quality and shape, or adversely affects the surface properties and chemical conversion treatment, especially the aesthetics of the steel sheet surface. Poor influence on stickiness It is an element that has an effect, and is limited to 1.0% or less in the present invention. If the Si content is 1.0% or less, the above-mentioned adverse effects can be suppressed. It is desirable that the content of Si be 0.5% or less, particularly for applications requiring beautiful surface of the coated steel sheet surface.
  • Mn is an effective element for preventing hot cracking due to S. It is preferable to add Mn in accordance with the amount of S contained.Mn has a great effect on refining crystal grains, and Mn is added. It is desirable to use it for improvement. From the viewpoint of stably fixing S, the content of Mn is desirably 0.1% or more. Mn is an element that increases the strength of the steel sheet, and when more strength is required, it is desirable to contain 0.5% or more. It is more preferably at least 0.8%.
  • Mn content When the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet to be sealed against the fluctuation of the hot rolling conditions, especially the variation of the strain aging hardening characteristics are remarkably improved.
  • Mn if Mn is excessively contained in excess of 2.0%, the detailed mechanism is unknown, but it tends to increase the hot deformation resistance, and also tends to deteriorate the weldability and the weldability.
  • Mn was limited to 2.0% or less because the formation of ferrite was remarkably suppressed, ductility was remarkably reduced, and the r-value tended to be remarkably reduced.
  • the content is preferably 1.5% or less.
  • P is a useful element as a solid solution strengthening element for steel, and is preferably contained in an amount of 0.002% or more from the viewpoint of increasing strength, and more preferably contained in an amount of 0.02% or more when high strength is required. Is more preferred. On the other hand, if it is contained excessively, it makes the steel brittle and further deteriorates the stretch flangeability of the steel sheet. In addition, P has a strong tendency to segregate in steel, resulting in embrittlement of the weld. Therefore, P was limited to 0.1% or less. In applications where stretch flangeability and weld toughness are particularly important, P is preferably set to 0.08% or less. More preferably, it is not more than 0.06%.
  • S 0.02% or less
  • S is an element that exists as an inclusion in the steel sheet and reduces the ductility of the steel sheet, and furthermore, deteriorates the corrosion resistance.It is preferable that S is reduced as much as possible.In the present invention, S is limited to 0.02% or less. . In particular, for applications requiring good workability, S is preferably set to 0.015% or less. When particularly excellent stretch flangeability is required, S is preferably set to 0.010% or less. Although the detailed mechanism is unknown, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening property of the steel sheet at a high level.
  • A1 is an element that acts as a deoxidizing agent, improves the cleanliness of steel, and further refines the structure of the steel sheet.
  • the content of A1 is preferably 0.001% or more.
  • N in a solid solution state is used as a strengthening element.
  • Aluminum-killed steel containing A1 within an appropriate range has better mechanical properties than conventional rimmed steel without A1 added. .
  • an excessive amount of A1 enhances the surface properties of the steel sheet and further significantly reduces the N in the solid solution state, making it difficult to obtain an extremely large amount of strain age hardening which is the main objective of the present invention. You. For this reason, in the present invention, A1 is limited to 0.02% or less.
  • A1 is 0.001 to 0.015%.
  • the present invention makes it possible to optimize this by optimizing the amount of other alloying elements and setting the annealing conditions in the optimum range. Is prevented.
  • N is an element that increases the strength of the steel sheet by solid solution strengthening and strain age hardening, and is the most important element in the present invention. Further, in the present invention, the cold rolling is performed by containing an appropriate amount of N, adjusting the A1 content to an appropriate value as described above, and further controlling the manufacturing conditions such as hot rolling conditions and annealing conditions. Ensure N in the solid solution state necessary and sufficient for the product or plating product. As a result, the effect of increasing the strength (yield stress and tensile strength) due to solid solution strengthening and strain age hardening is fully exhibited, with a tensile strength of 340 MPa or more and baking hardening amount.
  • N has an effect of lowering the transformation point, and it is effective to include N in the case of rolling of a thin material or the like in which rolling that does not want to greatly reduce the transformation point is desired.
  • N is less than 0.0050%, the above-described effect of increasing the strength is unlikely to appear stably.
  • N exceeds 0.0250%, the occurrence rate of internal defects in the steel sheet increases, and slab cracks and the like during continuous production become more frequent. For this reason, N was limited to the range of 0.0050 to 0.0250%.
  • N is preferably in the range of 0.0070 to 0.0200%, more preferably 0.0100 to 0.0170%. . If the N content is within the range of the present invention, there is no adverse effect on weldability and the like.
  • the amount of solute N is defined as a value obtained by subtracting the amount of precipitated N from the total amount of N in steel.
  • the method is determined by an electrolytic extraction analysis method using a constant potential electrolysis method as a result of comparative studies of various methods by the present inventors.
  • acid decomposition method, halogen method, and electrolysis method as a method of dissolving ground iron used for extraction analysis.
  • the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides.
  • the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed a good correspondence with the actual material change.
  • the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N.
  • the solid solution N content should be 0.0020% or more, In order to obtain a higher value, the content is preferably set to 0.0030% or more.
  • the upper limit of the amount of solute N is not particularly limited, but even if all the amount of N remains, the decrease in mechanical properties is small.
  • NZAl N content and A1 content ratio
  • NZA1 is preferably set to 0.6 or more from the viewpoint of stably improving the strain aging characteristics. More preferably, it is 0.8 or more.
  • Nb effectively acts to form an ash-ferrite phase in combination with B, and the present invention requires a content of 0.002% or more in the present invention.
  • the content exceeds 0.050%, the effect is saturated and the hot deformation resistance is significantly increased, so that hot rolling becomes difficult.
  • Nb was limited to the range of 0.002-0.050%. In addition, more preferably, it is 0.005 to 0.040%.
  • B is an element that effectively acts to form the ferrite phase in combination with Nb.
  • the content of 0.0001% or more is required.
  • the solute N which contributes to the strain age hardening characteristics is reduced. Therefore, B is limited to the range of 0.0001% to 0.0050%.
  • it is 0.0003-0.0030%. More preferably, it is 0.0005 to 0.0030%.
  • Group a 1.0% or less in total of one or more of Cu, Ni, Cr and Mo
  • Group b 0.1% or less in total of one or more of Ti and V
  • Group c One or two of Ca and REM in total 0.0010 to 0.010%
  • Group a elements are all elements that contribute to the increase in the strength of the steel sheet, and can be selected singly or in combination as necessary. Such an effect is recognized when Cu, Ni, Cr, and Mo are contained at 0.01% or more, respectively. However, if the content is too large, the hot deformation resistance increases, or the chemical conversion property and the surface treatment properties in a broad sense deteriorate, and the welded part is hardened and the formability of the welded part is deteriorated. Therefore, it is preferable that Cu, Ni, Cr, and Mo each alone be 1.0% or less, 1.0% or less, 0.5% or less, and 0.2% or less. Is preferably 1.0% or less in total.
  • Ti and V are elements that contribute to the refinement and uniformization of crystal grains, and can be selected as necessary and contained alone or in combination. Such an effect is recognized when the contents of Ti and V are each 0.005% or more. However, if the content is too large, the hot deformation resistance increases, and the chemical conversion property and the surface treatment properties in a broad sense deteriorate. In addition, there is an adverse effect of reducing solid solution N. Therefore, Ti and V alone are preferably 0.1% or less and 0.1% or less, respectively, and when they are contained in combination, the total content is preferably 0.1% or less.
  • Group c elements Ca and REM are both elements that are useful for controlling the morphology of inclusions, and if there is a requirement for stretch-flange formability, they are preferably contained alone or in combination. If the total of the elements in group d is less than 0.0010%, the effect of controlling the morphology of inclusions will be insufficient, whereas if it exceeds 0.010%, the occurrence of surface defects will be noticeable. For this reason, it is preferable to limit the total of the elements of the d group to the range of 0.0010 to 0.010%, whereby the stretch flangeability can be improved without the occurrence of surface defects.
  • the steel sheet of the present invention has a structure composed of a ferrite phase having an area ratio of 5% or more and a ferrite phase having an average crystal grain size of 20 / zm or less.
  • the cold-rolled steel sheet of the present invention contains 5% or more in area ratio of the ferrite-ferrite phase.
  • the presence of at least 5% of the ferrite phase provides good ductility and a large amount of strain age hardening.
  • the presence of the asymmetric ferrite phase causes the strain to be accumulated very effectively inside the pre-strain machining before aging.
  • the presence of the ferrite phase is effective in improving the deterioration of aging at room temperature and making it non-aging at room temperature.
  • the area ratio of the ferrite-ferrite phase is 10% or more.
  • the presence of a large amount of ash-ferrite phase exceeding 20% has a problem that the r-value decreases.
  • the area ratio of the fermentation phase is 5% or more, preferably 10% or more and 20% or less.
  • the ferrite phase referred to in the present invention is a low-temperature transformation phase unique to ultra-low carbon steel without a carbide therein, such as the composition of the present invention, which is mainly a normal ferrite phase obtained by observation with an optical microscope. Is a phase that is clearly identifiable, has a high internal dislocation density, and is harder than the polygonal ferrite phase.
  • the ferrite phase has the following features: (1) crystal grains in which the grain boundaries are irregularly angular, (2) crystal grains existing along the grain boundaries such as precipitates, and (3) scratch-like patterns.
  • One or a combination of grains (a number of sub-boundaries are found in relatively large second-phase grains). It can be clearly distinguished from Nalferite.
  • the corroded color tone in the grains is different from martensite and bainite, and is almost the same as ordinary polygonal ferrite, so that it can be clearly distinguished from martensite and bainite.
  • the dispersoid density in the ferrite phase is very high near the grain boundary and in the Z or intragranular phase. The lower part is in the form of a layer.
  • the cold-rolled steel sheet of the present invention is intended for a steel sheet for automobiles that requires high formability, and in order to ensure ductility, a phase other than the ferrite phase is a ferrite phase. If the area ratio of the ferrite phase is less than 80%, it is difficult to secure the ductility necessary for an automotive steel sheet requiring workability and a high r-value. If better ductility is required, the area ratio of the ferrite phase is preferably at least 80%, more preferably at least 85%.
  • the ferrite in the present invention refers to a so-called polygonal ferrite in which no distortion remains.
  • Average grain size of ferrite phase 20 / i in or less
  • the average crystal grain size a value calculated from the cross-sectional structure photograph by the quadrature method specified by ASTM and a nominal grain size determined by the cutting method also specified by ASTM (for example, Umemoto et al. (1984), 334).
  • ASTM for example, Umemoto et al. (1984), 334).
  • a predetermined amount of solute N is ensured at the product stage.
  • strain aging has occurred.
  • the curing characteristics varied, and it was found that one of the main factors was the crystal grain size.
  • a high BH content and ATS can be obtained stably by setting the average crystal grain size to at least not more than, preferably not more than 15 / zm.
  • the detailed mechanism is unknown, it is presumed to be related to the segregation and precipitation of alloying elements at the grain boundaries, as well as the effects of processing and thermal history on these.
  • the average crystal grain size of the ferrite phase is preferably 20 / zm or less, and more preferably 15 # ⁇ or less.
  • the cold-rolled steel sheet of the present invention having the above-described composition and structure has a tensile strength (TS) of 340 MPa or more and about 590 MPa or less, a high r value of r value of 1.2 or more, and excellent strain aging. It is a cold rolled steel sheet having hardening characteristics. Steel sheets with TS below 340MPa cannot be widely applied to structural members. In order to further expand the applicable range, it is desirable that T S be 400 MPa or more. If the r-value is less than 1.2, it cannot be applied to a wide range of press-formed parts. The preferred range of the r value is 1.3 or more.
  • the conventional paint baking condition is 170 ° C X 20 min as standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, Hardening is also achieved with (lower temperature) treatment, in other words, a wider range of aging conditions is possible. In general, in order to increase the amount of hardening, it is advantageous to hold at a higher temperature and for a longer time, unless softening is caused by excessive aging.
  • the lower limit of the heating temperature at which hardening becomes significant after pre-deformation is approximately 100 ° C.
  • the heating temperature exceeds 300, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous.
  • the holding time is about 30 s or more when the heating temperature is about 200 ° C, almost sufficient curing can be achieved. In order to obtain a still more stable deterioration, the holding time is preferably 60 s or more. However, holding for more than 20 min is not practical because it does not allow further curing and the production efficiency drops significantly.
  • the aging treatment conditions were evaluated at a heating temperature of 170 ° C. and a holding time of 20 min under the conventional paint baking treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention.
  • the method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used.
  • the strength of automotive components needs to be able to withstand complex external stress loads. Therefore, not only the strength characteristics in a small strain range but also the strength characteristics in a larger strain range are important for a material steel plate.
  • the present inventors set the BH amount (corresponding to the strength characteristic in a relatively small strain range) of the steel sheet of the present invention to be used as a material for automobile parts to 80 MPa or more, and set the ⁇ TS (Corresponding to the strength characteristics of the region) should be 40MPa or more. It is more preferable that the BH content be lOOMPa or more, and ⁇ TS 50MPa or more. Also, by setting the heating temperature during aging treatment to a higher temperature side and setting Z or the holding time to a longer time side, the BH amount and ATS can be further increased.
  • the steel sheet of the present invention preferably has a thickness of 3.2 mm or less. Further, in the present invention, there is no problem even if the surface of the above-mentioned cold rolled steel sheet of the present invention is electroplated or melted. These plated steel sheets also show the same amount of TS, BH and ATS as before plating.
  • any of electroplating, hot-dip galvanizing, alloyed hot-dip galvanizing, electro-tin plating, electro-chrome plating, and electro-nickel plating can be preferably applied.
  • a method for manufacturing a steel sheet according to the sixth invention will be described.
  • the steel sheet of the present invention is basically a steel slab having a composition in the above-mentioned range, which is subjected to rough rolling after heating to form a sheet par.
  • the sheet par is subjected to finish rolling.
  • the slab used in the production method of the present invention is desirably produced by a continuous production method in order to prevent macroscopic segregation of components, but may be produced by an ingot-making method or a thin slab production method.
  • direct-feed rolling in which the slab is placed in a heating furnace and rolled as it is without cooling, or slight heat retention is performed. Energy saving processes such as direct rolling, in which rolling is performed immediately afterwards, can be applied without any problems.
  • direct rolling is one of the useful techniques for effectively securing solid solution N.
  • Slab heating temperature 1000. C or more
  • the slab heating temperature is preferably set to 1000 ° C. or higher in order to secure a necessary and sufficient amount of solute N in the initial state and to satisfy a target value of the amount of solute N in the product.
  • the acid 1280 due to an increase in loss due to an increase in chemical weight. It is desirable to be C or less.
  • the slab heated under the conditions described above is converted into a sheet par by rough rolling.
  • the conditions for rough rolling do not need to be particularly defined, but may be determined according to a conventional method. However, from the viewpoint of securing the amount of dissolved N, it is desirable to carry out the reaction in as short a time as possible.
  • the sheet bar is finish-rolled into a hot-rolled sheet.
  • continuous sheet rolling is performed by joining the adjacent sheet pars between rough rolling and finish rolling.
  • the joining means it is preferable to use a laser welding method, an electron beam welding method, or the like even in a pressure welding method.
  • Continuous rolling eliminates the so-called rolling unsteady portions at the leading and trailing ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length of the coil (material to be processed). .
  • This is extremely effective in improving the cross-sectional shape and dimensions of not only hot-rolled steel sheets but also cold-rolled steel sheets. Further, even when cooling on a hot run table after rolling, tension can always be applied, so that the steel plate shape can be kept good.
  • a sheet par edge heater that heats the width direction end of the sheet par and a sheet par heater that heats the length direction end of the sheet par on the side of the finish rolling mill between the rough rolling and the finish rolling are provided. It is preferable to use one or both of them to make the temperature distribution in the width direction and the longitudinal direction of the sheet par uniform. As a result, it is possible to further reduce the material variation in the steel sheet. It is preferable that the sheet edge heater and the sheet heater are of an induction heating type.
  • a sheet-per-edge heater it is desirable to use a sheet-per-edge heater to compensate for the temperature difference in the width direction.
  • the amount of heating at this time depends on the steel composition and the like, but is preferably set so that the temperature distribution range in the width direction at the finish rolling exit side is approximately 20 ° C. or less.
  • sheet The temperature difference in the longitudinal direction is compensated by the per heater.
  • the heating amount at this time is preferably set so that the temperature at the end of the length is approximately 20 ° C higher than the temperature at the center.
  • Finishing rolling exit side temperature 800 ° C or more
  • the finish rolling exit side temperature FDT is 800 or more in order to obtain a uniformly fine hot-rolled base plate structure. If the FDT is below 800 ° C, the microstructure of the steel sheet becomes non-uniform, the processed microstructure remains in part, and the non-uniform microstructure remains after passing through the cold rolling annealing process. For this reason, there is an increased risk of various defects occurring during press forming. In addition, if a high winding temperature is used to avoid the remaining of the processed tissue, coarse crystals are generated, and the same problem occurs. In addition, when the winding temperature is set to a high temperature, a remarkable decrease in the amount of solute N occurs, so that it is difficult to obtain a target tensile strength of 340 MPa or more.
  • the finish rolling exit temperature FDT was set to 800 or more. In order to further improve the mechanical properties, it is desirable that FDT be 820 ° C or more. From the viewpoint of improving the r value, it is more preferable that FDT is equal to or higher than the Ac 3 transformation point. In particular, the upper limit of FDT is not specified, but if it is excessively high, scale flaws and the like will be remarkable. It is preferable that FDT is approximately 1000 or less.
  • Winding temperature 800 ° C or less
  • C T is preferably 800 or less. If the CT is less than 200, the shape of the steel sheet tends to be disturbed, and there is a high risk of causing problems in actual operation, and the uniformity of the material tends to decrease. For this reason, CT is desirably 200 ° C or more. When more uniform material is required, CT is preferably 300 or more. The value is more preferably 350 or more.
  • lubricating rolling may be performed in order to reduce a hot rolling load.
  • the coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10.
  • Hot rolling operation is stabilized.
  • the hot-rolled sheet that has been subjected to the above-described hot rolling step is then subjected to pickling and cold rolling in a cold-rolling step to become a cold-rolled sheet.
  • the conditions for pickling may be generally known conditions, and are not particularly limited. If the scale of the hot rolled sheet is extremely thin, cold rolling may be performed immediately without performing pickling.
  • the cold rolling conditions may be generally known conditions, and are not particularly limited. In addition, it is preferable that the cold rolling reduction is 60% or more from the viewpoint of ensuring the uniformity of the tissue.
  • the cold rolled sheet is then subjected to a cold rolled sheet annealing step consisting of continuous annealing and cooling.
  • Annealing at a temperature in the coexisting region of ferrite-austenite-to-ni phase forms a ferrite-ferrite phase.
  • a high r value is obtained because the (111) texture develops strongly in the ferrite phase.
  • the r-value decreases because the texture of the steel sheet is randomized by reverse transformation and transformation.
  • the annealing temperature of the continuous annealing is limited to a temperature not lower than the recrystallization temperature and in the ferrite to austenite dual phase coexisting region.
  • the temperature is preferably set so that the austenite fraction is 10% or more and 50% or less from the viewpoint of the stability of the r value. Further, if the continuous annealing temperature is lower than the recrystallization temperature, the ductility becomes low, so that it can be applied only to special applications limited to automotive parts.
  • the holding time of the continuous annealing time is preferably as short as possible from the viewpoints of production efficiency, refining the structure, and securing the amount of solute N.
  • the holding time is preferably at least 10 s from the viewpoint of operation stability, and is preferably at most 90 s from the viewpoint of microstructuring of the structure and securing the amount of dissolved N.
  • it is more preferable to set it to 20 s or more.
  • Cooling after continuous annealing Cooling at a cooling rate of 10 to 300 ° C / s to a temperature range of 500 ° C or less Cooling after soaking in continuation annealing is important from the viewpoint of microstructural refinement, formation of the ferrite-ferrite phase, and securing of the solute N content.
  • temper rolling or leveling may be performed for the purpose of shape correction and roughness adjustment. If the total elongation of the temper rolling or leveling is less than 0.5%, the intended purposes of shape correction and roughness adjustment cannot be achieved. On the other hand, if it exceeds 10%, the ductility is reduced. Note that the content is more preferably 5% or less from the viewpoint of ensuring ductility.
  • temper rolling and leveler processing are different, it has been confirmed that there is no significant difference between the two. Temper rolling and leveling are effective even after plating. The reason for limiting the composition of the high-tensile cold-rolled steel sheet of the seventh invention will be described.
  • C is an element that increases the strength of the steel sheet, and contains 0.025% or more in order to uniformly and finely control the structure, which is an important component of the present invention, and to secure a sufficient amount of martensite phase.
  • the carbide fraction in the steel sheet becomes excessive, and the ductility and the formability are significantly reduced.
  • the C content exceeds 0.15%, spot weldability, arc weldability, etc. will be significantly reduced.
  • C is limited to the range of 0.025 to 0.15%.
  • the content is preferably set to 0.08% or less. Further, for applications requiring particularly good ductility, the content is more preferably 0.05% or less.
  • Si is a useful element that can increase the strength of a steel sheet without significantly reducing the ductility of the steel, and is preferably contained at 0.005% or more, more preferably at 0.1% or more. .
  • Si significantly raises the transformation point during hot rolling, making it difficult to ensure quality and shape, or adversely affects the surface properties and chemical treatment, especially the beauty of the steel sheet surface.
  • It is an element that also has an adverse effect on plating, and is limited to 1.0% or less in the present invention.
  • the content of Si is 1.0% or less, the above-mentioned adverse effects can be suppressed to a low level.
  • the content of Si is desirably 0.5% or less.
  • Mn is an effective element for preventing hot cracking due to S. It is preferable to add Mn in accordance with the amount of S contained.Mn has a great effect on refining crystal grains, and Mn is added. It is desirable to use it for improvement. Furthermore, Mn is an extremely effective element for stably generating martensite during rapid cooling after continuous annealing. From the viewpoint of stably fixing S, the content of Mn is desirably 0.2% or more. Mn is an element that increases the strength of the steel sheet, and when a strength of more than 500 MPa in T S is required, it is preferable to contain 1.2% or more. It is more preferably at least 1.5%.
  • the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet with respect to the fluctuation of the hot rolling conditions, especially the variation of the strain aging hardening characteristics are remarkably improved.
  • Mn is excessively contained in excess of 2.0%, it becomes difficult to obtain a high r value, which is one of the important requirements of the present invention, and ductility is significantly reduced. Is limited to 2.0% or less.
  • the content is preferably 1.7% or less.
  • P is an element useful as a solid solution strengthening element for steel, and preferably contains 0.001% or more, more preferably 0.015% or more, from the viewpoint of increasing strength.
  • excessive This embrittles the steel and further deteriorates the stretch flangeability of the steel sheet.
  • p has a strong tendency to segregate in steel, which results in embrittlement of the weld. For this reason,
  • P was limited to 0.08% or less. In applications where stretch flangeability and weld toughness are particularly important, P is preferably set to 0.04% or less.
  • S is an element that exists as an inclusion in the steel sheet and reduces the ductility of the steel sheet, and furthermore, deteriorates the corrosion resistance.It is preferable that S is reduced as much as possible.In the present invention, S is limited to 0.02% or less. . In particular, for applications requiring good workability, S is preferably set to 0.015% or less. Further, when particularly excellent stretch flangeability is required, S is preferably 0.008% or less. Although the detailed mechanism is unknown, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening property of the steel sheet at a high level.
  • A1 is an element that acts as a deoxidizing agent, improves the cleanliness of steel, and further refines the structure of the steel sheet.
  • the content of A1 is preferably 0.001% or more.
  • N in a solid solution state is used as a strengthening element.
  • Aluminum-killed steel containing A1 within an appropriate range has better mechanical properties than conventional rimmed steel without A1 added. .
  • excessive A1 content deteriorates the surface properties of the steel sheet, and further significantly reduces N in the solid solution state, making it difficult to obtain an extremely large amount of strain age hardening. For these reasons, A1 is limited to 0.02% or less in the present invention.
  • A1 is preferably set to 0.001 to 0.015%.
  • the reduction of the A1 content may lead to coarsening of the crystal grains, but in the present invention, it is effective to limit the other alloying elements to the optimum amount and to set the annealing conditions in the optimum range. Has been prevented.
  • N is an element that increases the strength of the steel sheet by solid solution strengthening and strain age hardening, and is the most important element in the present invention.
  • an appropriate amount of N is contained.
  • the Al content is adjusted to an appropriate value and controlling the manufacturing conditions such as hot rolling conditions and annealing conditions, N in the solid solution state necessary and sufficient for cold rolled products or plated products can be obtained.
  • Secure As a result, the effect of increasing the strength (yield stress and tensile strength) by solid solution strengthening and strain age hardening is fully exhibited, with a tensile strength of 440 MPa or more, bake hardening amount (BH amount) of 80 MPa or more, and before and after strain aging treatment.
  • the target value of the mechanical properties of the steel sheet of the present invention that is, the increase amount of the tensile strength ⁇ TS of 40 MPa or more can be stably obtained.
  • N is less than 0.0050%, the above-described effect of increasing the strength is unlikely to appear stably.
  • N exceeds 0.0250%, the rate of partial cracking of the steel sheet will increase, and slab cracking will occur more frequently during continuous forming. For this reason, N was limited to the range of 0.0050 to 0.0250%. It is more preferable that N is in the range of 0.0070% to 0.0170% from the viewpoints of material stability and yield improvement in consideration of the entire manufacturing process. If the N content is within the range of the present invention, there is no adverse effect on weldability and the like.
  • the solid solution N amount is defined as a value obtained by subtracting the precipitated N amount from the total N amount in the steel.
  • the electrolytic extraction analysis method it is effective to obtain the amount by the electrolytic extraction analysis method using the potentiostatic electrolysis method as a result of comparative studies of various methods by the present inventors.
  • acid decomposition method, halogen method and electrolysis method as a method for dissolving base iron used for extraction analysis.
  • the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides. Electrolyte at a constant potential using an acetyl.acetone system as the electrolytic solution.
  • the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed a good correspondence with the actual material change.
  • the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N.
  • the amount of solute N is preferably 0.0020% or more.
  • the content is preferably 0.0030% or more.
  • the upper limit of the amount of solute N is not particularly limited, but even if all of the added N remains, the decrease in mechanical properties is small.
  • NZA1 (Ratio of N content and A1 content): 0.3 or more
  • Group d 1% or more of Cu, Ni, Cr, Mo, 1.0% or less in total
  • g group Ca or REM 1 or 2 kinds in total 0.0010 to 0.010.%
  • Group d elements are all elements that contribute to the increase in the strength of the steel sheet, and can be selected as necessary and contained alone or in combination. Such an effect is recognized when the content of Cu, Ni, Cr, and Mo is 0.005% or more, respectively. However, when the content is too large, the hot deformation resistance increases, or the chemical conversion property and the surface treatment properties in a broad sense deteriorate, and the welded part is hardened and the weldability is deteriorated. Also, the r value tends to decrease. For this reason, it is preferable that the total of the elements in group a be 1.0% or less. When Mo is contained in a large amount of 0.05% or more, the r value may be significantly reduced. In the present invention, when Mo is contained, the content is preferably limited to less than 0.05%.
  • Nb, Ti, and V are elements that contribute to the refinement and uniformity of crystal grains. Yes, and can be selected alone or in combination as needed. Such an effect is observed when Nb, Ti, and V are each contained at 0.005% or more. However, if the content is too large, the hot deformation resistance increases, and the chemical conversion property and the surface treatment properties in a broad sense deteriorate. For this reason, it is preferable that the total of the elements of group b be 0.1% or less.
  • B is an element that has the effect of improving the hardenability of steel, and increases the fraction of low-temperature transformation phases other than the ferrite phase, as necessary to increase the strength of steel. Can be contained. Such an effect is recognized when B is contained at 0.0005% or more. However, if the amount is too large, the hot deformability decreases, and BN is formed to reduce the solute N. Therefore, B is preferably set to 0.0030% or less.
  • Elements of group g are both elements that are useful for controlling the morphology of inclusions, and particularly when stretch flangeability is required, it is preferable to include them alone or in combination.
  • the total of the elements in group d is less than 0.0010%, the effect of controlling the morphology of inclusions is insufficient, while if it exceeds 0.010%, the occurrence of surface defects becomes noticeable.
  • it is preferable to limit the total number of elements in the d group to the range of 0.0010 to 0.010%, thereby improving the stretch flange workability without generating surface defects. it can.
  • the cold-rolled steel sheet according to the present invention is intended for a steel sheet for automobiles that requires a certain degree of workability, and has a structure containing 80% or more of a fly phase in area ratio in order to ensure ductility. If the area ratio of the ferrite phase is less than 80%, it will be difficult to secure the required ductility as an automotive steel sheet that requires workability. If even better ductility is required, the area ratio of the ferrite phase should be 85% or more.
  • the ferrite according to the present invention refers to a so-called polygonal ferrite in which no distortion remains.
  • Average grain size of ferrite phase 10 / x m or less
  • the average crystal grain size is calculated by the quadrature
  • the larger of the value calculated by the above and the nominal particle size (see, for example, Umemoto et al .: Heat treatment, 24 (1984), 334) determined by the cutting method also specified in ASTM is adopted.
  • a predetermined amount of solute N is ensured at the product stage.
  • strain aging has occurred.
  • the hardening characteristics may flicker, and it has been found that one of the main factors is the crystal grain size.
  • the average grain size is at least 10 / z ra or less, preferably 8 / ⁇ or less, a stable high BH content and ATS can be obtained.
  • the detailed mechanism is unknown, it is presumed to be related to the bias and precipitation of alloying elements at the grain boundaries, and the effects of processing and thermal history on these.
  • the average crystal grain size of the ferrite phase needs to be ⁇ ⁇ ⁇ ⁇ or less, preferably 8 / ⁇ or less.
  • the present invention provides a structure containing 80% or more of ferrite having an average crystal grain size of 10 / Di or less by area ratio. I do.
  • Martensite phase area ratio 2% or more
  • the cold-rolled steel sheet of the present invention contains, as the second phase, a martensite phase in an area ratio of 2% or more.
  • a martensite phase in an area ratio of 2% or more.
  • the area ratio of the martensite phase is preferably 5 ° / 0 or more.
  • the presence of a large amount of martensite phase exceeding 20% has a problem that ductility is reduced. Therefore, the area ratio of the martensite phase is 2% or more, preferably 5% or more and 20% or less.
  • the ferrite phase As the second phase, there is no problem that pearlite, bainite and residual austenite exist in addition to the above-mentioned martensite phase, but in the present invention, the ferrite phase The fraction must be 80% or more and the martensite phase fraction must be 2% or more. The total area ratio of perlite, payite and residual austenite is limited to less than 18%.
  • the cold-rolled steel sheet of the present invention having the above-described composition and structure has a tensile strength (TS) of 440 MPa or more and about 780 MPa or less, and further has a high r-value of 1.2 or more by controlling the texture of the matrix ferrite. It is a cold-rolled steel sheet that has an r-value and excellent strain aging hardening characteristics. Steel sheets with TS below 440MPa cannot be widely applied to members with structural elements. To further expand the application range, it is desirable that T S be 500 MPa or more. If the r value is less than 1.2, it cannot be applied to a wide range of press-formed parts. The preferred range of the r value is 1.4 or more.
  • excellent strain aging hardening characteristics means that, as described above, after pre-deformation with a tensile strain of 5%, when subjected to aging treatment at a temperature of 170 for 20 minutes, the deformation stress before and after this aging treatment is increased.
  • the amount of prestrain (prestrain) is an important factor.
  • the present inventors have investigated the effect of the amount of pre-strain on the strain aging hardening characteristics, assuming a deformation mode applied to a steel sheet for automobiles. As a result, (1) the deformation stress in the above-mentioned deformation mode is extremely low. Except in the case of deep drawing, it can be roughly organized by the strain (tensile strain) equivalent to one axis. (2) In actual parts, the strain equivalent to one axis is more than about 5%. It has been found that the strength (YS and TS) obtained after the strain aging treatment of% corresponds well. Based on this finding, in the present invention, the pre-strain of the strain aging treatment was determined to be 5% tensile strain.
  • the conventional paint baking condition is 170 ° C X 20 min as standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, hardening is achieved even with milder (lower temperature) treatment, in other words, aging conditions can be broadened. And it is possible. In general, in order to increase the amount of hardening, it is advantageous to hold at a higher temperature and for a longer time, unless softening is caused by excessive aging.
  • the lower limit of the heating temperature at which hardening is remarkable after pre-deformation is approximately 100.
  • the heating temperature exceeds 300, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous.
  • the holding time is about 30 s or more when the heating temperature is about 200, almost sufficient curing can be achieved.
  • the holding time is preferably 60 s or more. However, holding for more than 20 min is not practical because it does not allow further hardening and significantly reduces power production efficiency.
  • the holding time was evaluated at 170 min, which is the heating temperature under the conventional paint baking treatment conditions, and the holding time was evaluated at 20 min as the aging treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention.
  • the method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used.
  • the strength of automotive components needs to be able to withstand complex external stress loads, so that not only the strength characteristics in a small strain range but also the strength characteristics in a large strain range are important for a material steel plate.
  • the present inventors set the BH amount of the steel sheet of the present invention, which is to be used as a material for automobile parts, to be 80 MPa or more and the ⁇ TS amount to be 40 MPa or more. More preferably, the amount of BH should be lOOMPa or more, and ⁇ TS should be 50MPa or more.
  • the heating temperature during the aging treatment to a higher temperature side and / or the holding time to a longer time side, the BH amount and the ATS amount can be further increased.
  • the steel sheet of the present invention has an advantage that it can be expected to increase the strength by about 40% of the full aging simply by leaving it at room temperature for about one week without heating after forming.
  • the steel sheet of the present invention even if left untreated at room temperature for a long time, It also has the advantage that conventional aging steel sheets do not have deterioration effect (phenomenon of increasing YS and decreasing El (elongation)), which is not possible with conventional aging steel sheets.
  • aging at room temperature for 3 months before press forming increases the YS by 3 OMPa or less, decreases the elongation by 2% or less, and increases the yield point elongation. Recovery must be less than 0.2%.
  • the steel sheet of the present invention is basically a steel slab having a composition in the above-mentioned range, which is subjected to rough rolling after heating to form a sheet par.
  • the sheet par is subjected to finish rolling.
  • the slab used in the production method of the present invention is desirably produced by a continuous production method in order to prevent macroscopic segregation of components, but may be produced by an ingot-making method or a thin slab production method.
  • direct-feed rolling in which the slab is placed in a heating furnace and rolled as it is without cooling, or slight heat retention is performed. Energy saving processes such as direct rolling, in which rolling is performed immediately afterwards, can be applied without any problems. In particular, direct rolling is one of the most useful technologies for effectively securing solid solution N.
  • the slab heating temperature is preferably at least 1000 ° C in order to secure a necessary and sufficient amount of solute N in the initial state of hot rolling and to satisfy the target value of solute N in the product. It is desirable to set the temperature to 1280 ° C or less because of the increase in loss due to the increase in oxidation weight.
  • the slab heated under the conditions described above is converted into a sheet par by rough rolling.
  • the conditions for rough rolling do not need to be particularly defined, but may be determined according to a conventional method. However, from the viewpoint of securing the amount of dissolved N, it is desirable to carry out the reaction in as short a time as possible.
  • the sheet par is finish-rolled into a hot-rolled sheet.
  • continuous sheet rolling is performed by joining the adjacent sheet pars between the rough rolling and the finish rolling.
  • a joining means it is preferable to use a laser welding method, an electron beam welding method, or the like even in a pressure welding method.
  • Continuous rolling eliminates the so-called unsteady rolling part at the front and rear ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length and width of the coil (material to be processed). .
  • This is extremely effective in improving the cross-sectional shape and dimensions of not only hot-rolled steel sheets but also cold-rolled steel sheets. Further, even when cooling on a hot run table after rolling, tension can always be applied, so that the steel plate shape can be kept good.
  • a sheet par edge heater for heating the width end of the sheet par and a sheet par heater for heating the length end of the sheet par are provided. It is preferable to use both of them to equalize the temperature distribution in the width direction of the sheet par. Thereby, the variation in the material in the steel sheet can be further reduced.
  • the sheet-per-edge heater and the sheet-bar heater are preferably of an induction heating type. It is desirable to use a sheet-per-edge heater to compensate for the temperature difference in the width direction.
  • the amount of heating at this time depends on the steel composition and the like, but is preferably set so that the temperature distribution range in the width direction at the finish rolling exit side is approximately 20 ° C. or less.
  • the temperature difference in the longitudinal direction is compensated for by the sheet heater. It is preferable that the heating amount in this case is set so that the temperature at the end of the length is approximately 20 ° C higher than the temperature at the center.
  • Finishing rolling exit temperature 800 or more
  • the finish-rolling exit temperature FDT should be 800 ° C or higher in order to obtain a uniform and fine hot-rolled base plate structure. If the FDT is lower than 800 ° C, the structure of the steel sheet becomes non-uniform, a part of the processed structure remains, and after the cold rolling annealing process, the non-uniform structure of the steel remains without disappearing. For this reason, there is a great risk that various problems occur during press forming. In addition, if a high winding temperature is used to avoid the remaining of the processed tissue, coarse crystals are generated, and the same problem occurs.
  • the finish rolling exit temperature FDT was set to 800 ° C or higher.
  • the FDT be 820 ° C or higher.
  • the upper limit of FDT is not specified, but if it is excessively high, scale flaws and the like will be noticeable.
  • FDT is generally preferably up to about 1000 e C.
  • the cooling after finish rolling is not particularly strictly limited, but the following conditions are desirable from the viewpoint of the material uniformity in the longitudinal and width directions of the steel sheet. That is, in the present invention, it is desirable that cooling be started immediately (within 0.5 seconds) after finishing rolling and the average cooling rate during cooling be 40 s or more. By satisfying this condition, the high-temperature region where A1N precipitates can be rapidly cooled, and N in solid solution can be secured effectively. If the cooling start time or cooling rate does not satisfy the above conditions, the grain growth will proceed too much, making it difficult to achieve a fine grain size, and the precipitation of A1N due to the strain energy introduced during rolling will be promoted. There is a possibility that the amount of solute N may be deficient, and the tissue tends to be uneven. From the viewpoint of ensuring uniformity of the material and the shape, the cooling rate is preferably suppressed to 300 eCZs or less. Winding temperature: 800 ° C or less
  • CT As the winding temperature CT decreases, the steel sheet strength tends to increase. In order to secure the target tensile strength T S 440 MPa or more, it is preferable that C T be 800 ° C. or less. If the CT is less than 200, the shape of the steel sheet is likely to be disturbed, and there is a high risk of causing a problem in actual operation, and the uniformity of the material tends to decrease. Therefore, it is desirable that CT is 200 or more. When more uniformity of the material is required, it is preferable that CT is 300 t or more. The temperature is more preferably 350 ° C. or higher. In the present invention, in finish rolling, lubricating rolling may be performed in order to reduce the hot rolling load.
  • the coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10.
  • the combination of lubrication rolling and continuous rolling further stabilizes the operation of hot rolling.
  • the hot-rolled sheet that has been subjected to the above-mentioned hot rolling step is then subjected to pickling and cold rolling in a cold-rolling step to become a cold-rolled sheet.
  • the conditions for pickling may be generally known conditions, and are not particularly limited. If the scale of the hot rolled sheet is extremely thin, cold rolling may be performed immediately without performing pickling.
  • the cold rolling conditions may be generally known conditions, and are not particularly limited. It is preferable that the cold rolling reduction is 40% or more from the viewpoint of ensuring the uniformity of the tissue. Next, the reasons for limiting the conditions of the cold rolling process will be described.
  • the cold-rolled sheet is then subjected to a cold-rolled sheet annealing process including box annealing and continuous annealing.
  • Box annealing temperature not less than recrystallization temperature and not more than 800
  • box annealing is performed on the cold-rolled sheet to control the texture of the ferrite phase as a base.
  • the r-value of the product plate can be increased.
  • This box annealing facilitates the formation of a (11 1) texture that is desirable for increasing the r-value on the product sheet.
  • the box annealing is preferably performed in an annealing atmosphere mainly containing nitrogen gas and containing 3 to 5% of hydrogen gas.
  • the heating and cooling rates may be the same as those in ordinary box annealing, and are generally about 30 °. It is about C Ar. Further, by using an annealing atmosphere gas of 100% hydrogen gas, a higher heating / cooling rate may be obtained.
  • Continuous annealing temperature Ac 1 transformation point or more (Ac 3 transformation point-20 ° C) or less
  • the continuous annealing temperature is not less than the Ac 1 transformation point and not more than the Ac 3 transformation point. Further, the holding time of the continuous annealing time is preferably as short as possible from the viewpoints of production efficiency, refining the structure, and securing the amount of solute N.
  • the holding time is preferably 10 s or more from the viewpoint of operation stability, and is preferably 120 s or less from the viewpoint of refining the structure and securing the amount of dissolved N.
  • Cooling after continuous annealing Cooling down to a temperature range of 500 ° C or less at a cooling rate of 10 to 300 Cooling after soaking in continuous annealing reduces the size of the structure, forms martensite, and secures the amount of solid solution N. It is important from the point of view.
  • continuous cooling is performed at a cooling rate of 1 Os or more to at least a temperature range of 500 ° C or less. If the cooling rate is less than 10 Vs, the required amount of martensite, a uniform and fine structure, and a sufficient amount of solute N cannot be obtained.
  • Overaging treatment condition After cooling after continuous annealing, residence time is 20 s or more in the temperature range of 350 V or less, which is lower than the cooling stop temperature of the cooling. Subsequent to the cooling stop after the soaking in the continuous annealing, an overaging treatment with a residence time of 20 s or more may be performed in a temperature range of 350 ° C or lower below the cooling stop temperature. By performing the overaging treatment, the amount of solute C can be selectively reduced while maintaining the amount of solute N. If the residence temperature range is lower than 350, it takes a long time to reduce the solid solution C, which leads to a decrease in productivity. Therefore, the temperature range is preferably 350 ° C or higher.
  • the residence time is preferably 120 s or less.
  • continuous annealing following box annealing is performed in a continuous hot-dip line, and continuous annealing is performed.
  • hot-dip galvanizing or further alloying can be performed to produce a hot-dip galvanized steel sheet.
  • Temper rolling or leveling elongation 0.2 to 15%
  • temper rolling or leveling may be performed for the purpose of shape correction and roughness adjustment. If the total elongation of the temper rolling or leveling is less than 0.2%, the intended purposes of shape correction and roughness adjustment cannot be achieved. On the other hand, if it exceeds 15%, the ductility is significantly reduced. Although the form of temper rolling and leveler processing are different, it has been confirmed that there is no significant difference between the two. Temper rolling and leveling are effective even after plating.
  • forming such as press forming
  • the strain introduced by press working is several percent to several tens of percent.
  • the amount of distortion varies depending on the molded parts, about 5 to 10% of distortion is introduced into the inner plates and structural members in the automotive field.
  • these formed parts are subjected to a heat treatment such as a paint baking treatment.
  • a heat treatment such as a paint baking treatment.
  • the strength of the formed parts can be effectively increased after the heat treatment.
  • a tensile test piece of JIS No. 5 size was sampled in the rolling direction, and 10% tensile strain was applied by a tensile tester. After that, heat treatment is performed, and then the tensile test is performed again.
  • heat treatment conditions should be 120 and 20 minutes. This test evaluates the properties of the completed parts that have been subjected to heat treatment following press forming.
  • the difference (ATS) between the tensile strength after such a tensile strain imparting-heat treatment and the tensile strength of the product is defined as the strength increasing heat treatment ability.
  • the amount of distortion introduced by molding is large or the heat treatment temperature after working is high.
  • the steel sheet of the present invention has sufficient strength even when the heat treatment temperature after forming is lower than before, that is, even when the heat treatment temperature is 200 ° C or less, when the applied strain amount is about 5 to 10% described above. Can be increased. Nevertheless, if the heat treatment temperature is lower than 120 ° C, a sufficient strength increasing effect cannot be obtained when the strain is low. On the other hand, when the heat treatment temperature after molding exceeds 350 eC , softening proceeds. Therefore, it is preferable that the heat treatment temperature after molding be about 120-350.
  • a heating method a method such as hot air heating, infrared furnace heating, hot bath heat treatment, electric current heating, and high frequency heating can be applied, and is not particularly specified. Alternatively, only the portion where the strength is to be increased may be selectively heated.
  • Example in the following examples the amount of solid solution N, microstructure, tensile properties, r value measurement, strain age hardening properties, and aging properties were investigated. The survey method is as follows.
  • the amount of solute N was determined by subtracting the amount of precipitated N from the total amount of N in the steel determined by chemical analysis.
  • the amount of precipitated N was determined by an analytical method using the above-described potentiostatic electrolysis method.
  • a specimen was taken from each cold-rolled annealed plate, and the microstructure of the cross section (C cross section) perpendicular to the rolling direction was imaged using an optical microscope or a scanning electron microscope, and the microstructure of ferrite was obtained using an image analyzer. The fraction, the type of the second phase and the tissue fraction were determined.
  • crystal grain size a value calculated by a quadrature method specified in ASTM from a cross-sectional structure photograph and a nominal particle size determined by a cutting method specified in ASTM from a cross-sectional structure photograph (for example, Umemoto et al .: Heat treatment, 24 (1994), 334) was used, whichever was greater.
  • YS 5% is the deformation stress when the product plate is pre-deformed 5%
  • YS BH and T SBH are the yield stress and tensile strength after pre-deformation-paint baking
  • TS is the product plate Is the tensile strength.
  • JIS No. 5 test pieces were sampled from the rolling direction (L direction), 45 ° direction (D direction) with respect to the rolling direction, and 90 ° direction (C direction) with respect to the rolling direction of each cold rolled annealed sheet.
  • L direction rolling direction
  • D direction 45 ° direction
  • C direction 90 ° direction
  • wo and t O are the width and thickness of the test piece before the test, and w and .t are the width and thickness of the test piece after the test.
  • the mean r value r mean was determined by
  • r L is the r value in the rolling direction (L direction)
  • r D is the r value in the 45 ° direction (D direction) with respect to the rolling direction (L direction)
  • rc is the rolling direction ( R value in the 90 ° direction (C direction) with respect to the L direction).
  • the calculation was performed based on changes in the elongation strain and the strain in the width direction.
  • JIS No. 5 test pieces were collected from each cold-rolled annealed sheet, and the test pieces were subjected to an aging treatment of 50 ⁇ X2001 ⁇ and then subjected to a tensile test. From the obtained results, the yield-elongation difference ⁇ -El before and after the aging treatment was determined, and the aging characteristics at room temperature were evaluated. If ⁇ -E1 was zero, it was evaluated as non-aging and excellent in normal temperature aging resistance.
  • the tensile strength after forming and heat treatment is the temperature at 120 ° C and the heat treatment temperature equivalent to conventional paint baking after taking a JIS No. 5 test piece from the product plate in the rolling direction and applying a 10% pre-strain. Heat treatment was performed at 170 ° C for 20 minutes, and the tensile strength was measured and determined.
  • the amount of decrease in total elongation due to normal temperature aging was calculated using the total elongation measured by taking a JIS No. 5 test piece from the product plate in the rolling direction and the JIS No. 5 test piece separately taken in the rolling direction. The difference from the total elongation measured after accelerating the heat aging treatment (holding at 100 ° C for 8 hours) was obtained.
  • a steel slab having the composition shown in Table 1 was converted into a hot-rolled strip with a thickness of 3.5 and then a cold-rolled strip with a thickness of 0.7 rain under the conditions shown in Table 2.
  • Table 3 shows the results of an investigation on the tensile strength and the r-value of the cold-rolled steel sheet and the alloyed hot-dip galvanized steel sheet thus obtained, and the change in the tensile strength after the forming and heat treatment.
  • both the cold-rolled steel sheet and the alloyed hot-dip galvanized steel sheet obtained according to the present invention have higher r-values and better strain age hardening characteristics than the comparative examples.
  • those having a crystal grain size of 20 / z ra or less have a small decrease in elongation due to aging at room temperature of 2.0% or less at ⁇ 1.
  • tensile strength TS 365 MPa
  • r value 1.7
  • a steel slab having the composition shown in Table 6 was hot-rolled under the conditions shown in Table 7 to obtain a hot-rolled sheet with a thickness of 3.5 mra. These hot-rolled sheets were cold-rolled under the conditions shown in Table 7 to obtain 0.7 mm-thick cold-rolled sheets, and these cold-rolled sheets were recrystallized and annealed under the conditions shown in the same table. Further, hot-dip galvanizing or alloyed hot-dip galvanizing was performed under the conditions shown in the same table. The obtained product plate was examined for the amount of solute N, microstructure, tensile properties, and strain age hardening properties. Table 8 shows the results.
  • the steel sheet according to the present invention has a TSX r value of ⁇ 750 MPa (for those in which B and one or more of Nb, Ti, and V are added in combination, a TSX r value of ⁇ 850 MPa) and a BH ⁇ 80 MPa ATS ⁇ 40 MPa, but one or more of these three characteristics do not reach the level of the present invention in the comparative example.
  • Molten steel with the composition shown in Table 9 was smelted in a converter and made into a steel slab by continuous casting. These steel slabs were heated under the conditions shown in Table 10, rough-rolled to form sheet sheets, and then hot-rolled by a hot rolling step of finish rolling under the conditions shown in Table 10.
  • the Ar 3 transformation point was measured using a working transformation measuring device (manufactured by Fuji Denki Co., Ltd.) under conditions simulating the hot finish rolling conditions, and is shown in Table 10.
  • These hot-rolled sheets were formed into cold-rolled sheets by a cold rolling process including pickling and cold rolling under the conditions shown in Table 10. Subsequently, continuous annealing was performed on these cold-rolled sheets under the conditions shown in Table 10 below. For some, temper rolling was performed after the cold rolling annealing process.
  • All of the examples of the present invention exhibit excellent ductility, a remarkably high BH content and ATS, excellent strain aging hardening properties, a high r-value with an average r-value of 1.2 or more, and non-aging properties at room temperature aging. It has excellent room temperature aging resistance.
  • the properties of galvanized steel sheets with hot-dip galvanized No. 4 and No. 10 steel sheets were 0.2 times the average r-value compared to cold-rolled steel sheets and 0.2% due to the restraint of width reduction of the coating layer. ⁇ E 1 decreased the properties by about 1%, but the strain age hardening property and the room temperature aging resistance were almost the same as the properties before plating.
  • the comparative examples out of the scope of the present invention do not have all of the target properties and have sufficient properties, depending on whether the ductility is deteriorated, the amount of BH, the ATS is small, or the aging deterioration is remarkable. It cannot be said that the steel sheet has
  • N was out of the preferred range of the present invention, solute N was small, and strain age hardening characteristics Is declining.
  • the hot rolling conditions and the cold rolling annealing conditions were out of the preferred ranges, and the microstructure was out of the range of the present invention. The aging resistance has deteriorated.
  • a steel having the composition shown in Table 12 was formed into a slab in the same manner as in Example 4, and the slab was heated under the conditions shown in Table 13 and roughly rolled to obtain a 25-long steel sheet. Then, the hot rolled sheet was formed by a hot rolling process in which finish rolling was performed under the conditions shown in Table 13. After rough rolling, the successive sheet pars on the entry side to the finish rolling were joined by a fusion welding method and continuously rolled. In addition, the width of the sheet par and the end in the length direction of the sheet par were controlled by using an induction heating type sheet par edge heater and a sheet par heater.
  • These hot-rolled sheets were pickled by a cold rolling process comprising cold rolling under the conditions shown in Table 13 to form cold-rolled sheets having a thickness of 1.6 mm.
  • a cold rolling process comprising cold rolling under the conditions shown in Table 13 to form cold-rolled sheets having a thickness of 1.6 mm.
  • these cold-rolled sheets were subjected to continuous annealing under the conditions shown in Table 13.
  • Example 4 With respect to the obtained cold rolled annealed sheet, the amount of dissolved N, microstructure, tensile properties, r value measurement, and strain age hardening properties were examined in the same manner as in Example 4. In addition, the tensile properties of each cold-rolled annealed sheet were examined at ten locations in the width direction and the longitudinal direction, and variations in yield strength, tensile strength, and elongation were examined.
  • each of the examples of the present invention had excellent strain age hardening characteristics and a high r-value, and showed a stably remarkably high BH content, ATS, and an average r-value despite fluctuations in production conditions.
  • the thickness accuracy and the shape of the product steel sheet were improved by performing continuous rolling and temperature adjustment in the longitudinal direction and the width direction of the sheet par, and the material variation was reduced to 1/2. did.
  • the elongation of the temper rolling was changed from 0.5 to 2% and the elongation of the leveler was changed from 0 to 1%, but there was no decrease in strain age hardening characteristics.
  • Molten steel with the composition shown in Table 15 was smelted in a converter, and was made into a steel slab by continuous casting. These steel slabs were heated under the conditions shown in Table 16 (with some hot flakes charged), rough-rolled to form a sheet pallet, and then subjected to hot rolling in a hot rolling process in which finish rolling was performed under the conditions shown in Table 16 It was a plate. In addition, in some sheet pars, successive sheet pars were joined to each other by a melt pressure welding method, and continuous rolling was performed.
  • All of the examples of the present invention exhibit excellent ductility, a remarkably high BH content and ATS, excellent strain aging hardening properties, a high r value of an average r value of 1.2 or more, and non-aging at room temperature. are doing.
  • the properties of the hot-dip galvanized steel sheets No. 17 and No. 18 shown in Table 17 were almost the same as those of the continuously annealed cold-rolled steel sheets.
  • the comparative examples out of the scope of the present invention do not have all of the target characteristics and have sufficient characteristics depending on whether the ductility is deteriorated, the BH amount and the ATS are small, or the aging deterioration is remarkable. It cannot be said that the steel sheet has
  • the slab heating temperature and the FDT were out of the preferred ranges of the present invention, the amount of solid solution N and the amount of martensite were out of the range of the present invention, and the average crystal grain size of the fiber was increased in the range of the present invention.
  • the r value, BH amount, and ATS have decreased.
  • the winding temperature after hot rolling was out of the range of the present invention, the amount of solute N departed from the range of the present invention, and the average crystal grain size of ferrite increased in the range of the present invention.
  • the r value, BH amount, and ATS have decreased.
  • the continuous annealing temperature was out of the preferred range of the present invention, no martensite was formed, and the average grain size of ferrite was out of the range of the present invention, so that the BH content and the TS decreased. ⁇ ⁇ -El is increasing.
  • no box annealing was performed, and the desired texture was not developed, so the r-value was particularly low.
  • the average particle size of ferrite and the area ratio of martensite are also out of the range of the present invention.
  • a steel having the composition shown in Table 18 was formed into a slab in the same manner as in Example 1, and the slab was heated under the conditions shown in Table 19 and roughly rolled to form a 30 mm thick sheet par.
  • a hot rolled sheet was obtained by a hot rolling step of performing finish rolling under the following conditions.
  • a part of the sheet bars, which were adjacent to each other on the entry side of the finish rolling after the rough rolling were joined by a melt pressure welding method and were continuously rolled.
  • the width of the sheet par and the end in the length direction of the sheet par were controlled by using an induction heating type sheet par edge heater and a sheet par heater.
  • hot-rolled sheets were formed into 1.6 ram-thick cold-rolled sheets by a cold rolling process including pickling and cold rolling under the conditions shown in Table 19.
  • box annealing was performed on these cold-rolled sheets under the conditions shown in Table 19, and then continuous annealing was performed in a continuous annealing furnace. Note that the annealing temperatures of the box annealing were all higher than the recrystallization temperature.
  • Each of the examples of the present invention had excellent strain age hardening characteristics and a high r-value, and showed a stably remarkably high BH content, ATS, and an average r-value despite fluctuations in production conditions. Further, in the present invention example, it was confirmed that by performing the continuous rolling and the temperature adjustment in the longitudinal direction and the width direction of the sheet par, the thickness accuracy and shape of the product steel sheet were improved, and the variation in the material was reduced.
  • the present invention it is possible to obtain a cold-rolled steel sheet in which TS is greatly increased by press forming-heat treatment while securing excellent deep drawability during press forming. From this cold-rolled steel sheet, there is an excellent effect that an electro-galvanized steel sheet, a hot-dip galvanized steel sheet, and an alloyed hot-dip galvanized steel sheet can be industrially manufactured.
  • Finish rolling end temperature is lower than Ar 3 transformation point.
  • the cooling conditions after finish rolling indicate the cooling start time (s) and the cooling rate (3 ⁇ 4 / s),
  • 2S000 ⁇ 3 2S000 OIOO0 0800, 0 U0 ⁇ 0210 TO0 iOO ⁇ TOO 00 800 00 18 00 81 00 9 ⁇ 00 00 ⁇ 0 60000 OiOO '0 -s s ⁇ 0 100 ⁇ 0 TOO ⁇ 0 ⁇ ⁇ 28 ⁇ 0 LI ⁇ 9200 ⁇ 0 d
  • Rate Particle size AF YS TS E 1 Average ⁇ YS ⁇ TS ⁇ E 1

Abstract

A cold rolled steel sheet excellent in strain aging hardening property, characterized in that it has a chemical composition, in mass %; C: less than 0.01 %, Si: 0.005 to 1.0 %, Mn: 0.01 to 1.0 %, Nb: 0.005 to 0.050 %, Al: 0.005 to 0.030 %, N: 0.005 to 0.040 %, B: 0.0005 to 0.0015 %, P: 0.05 % or less, S; 0.01 % or less and balance: substantially Fe, with the proviso that the composition satisfies the following formulae (1) and (2): N % ≥ 0.0015 + 14/93 • Nb% + 14/27 • Al% + 14/11 • B% ---(1) C % ≤ 12/93 • Nb% ---(2); and an alloyed hot dip galvanized steel sheet produced from the cold rolled sheet. The steel sheet exhibits a significantly enhanced tensile strength when it is subjected to press forming heat treatment, while maintaining excellent deep drawing properties in press forming.

Description

明 細 書  Specification
歪時効硬化特性を有する冷延鋼板および亜鉛めつき鋼板ならぴにそれらの製造方法 技術分野 '  Cold-rolled steel sheets and zinc-plated steel sheets having strain-age hardening properties and their manufacturing methods
この発明は、 建設部材、 機械構造用部品および自動車の構造用部品等、 構造上の 強度とくに変形時の強度および/または剛性が必要とされる箇所に用いられ、 プレ スなどによる加工成形後に強度上昇熱処理が施される成形体の素材鋼板として好適 な、 歪み時効硬化特性に優れた冷延鋼板、 電気亜鉛めつき鋼板、 溶融亜鉛めつき鋼 板、 および合金化溶融亜鉛めつき鋼板ならびにそれらの製造方法に関するものであ る。  INDUSTRIAL APPLICABILITY The present invention is used for parts where structural strength, especially strength and / or rigidity at the time of deformation is required, such as construction members, mechanical structural parts, and structural parts of automobiles. Cold rolled steel sheet, electro-zinc-plated steel sheet, hot-dip galvanized steel sheet, and alloyed hot-dip galvanized steel sheet, which are suitable as a material steel sheet of a molded body subjected to ascent heat treatment and have excellent strain age hardening characteristics It concerns the manufacturing method.
また、 本発明において、 「歪時効硬化特性に優れた」 とは、 引張歪 5 %の予変形 後、 170 の温度に 20rain 保持する条件で時効処理したとき、 この時効処理前後の 変形応力増加量 （B H量と記す； B H量 =時効処理後の降伏応力一時効処理前の予 変形応力） が 80MPa 以上であり、 かつ歪時効処理 （前記予変形 +前記時効処理） 前 後の引張強さ増加量 （A T Sと記す； A T S =時効処理後の引張強さ一予変形前の 引張強さ） が 40MPa 以上であることを意味する。 背景技術  Further, in the present invention, “excellent in strain age hardening characteristics” means that the amount of increase in deformation stress before and after this aging treatment when pre-deformation of 5% tensile strain and aging treatment at a temperature of 170 for 20 rains (BH amount; BH amount = pre-deformation stress before yield stress after aging treatment) is 80MPa or more, and tensile strength increase before and after strain aging treatment (pre-deformation + aging treatment) The amount (denoted as ATS; ATS = tensile strength after aging treatment-tensile strength before pre-deformation) is 40 MPa or more. Background art
薄鋼板のプレス成形体の製造に際しては、 プレス成形前は軟質としてプレス成形 'を容易にしておき、プレス成形後に硬化させて部品強度を高める方法として、 200 。C 未満で塗装焼付する方法があり、 かような塗装焼付用の鋼板として B H鋼板が開発 された。  In the production of press-formed thin steel sheets, the method of softening before press-forming is to make the press-forming 'easy, and after press-forming, it is hardened to increase the part strength. There is a method of baking paint below C, and a BH steel plate was developed as a steel plate for such baking paint.
例えば、 特開昭 55— 141526号公報には、 鋼中の C， N， A1含有量に応じて Nbを添 加し、 at %で Nb/ (固溶 C +固溶 N ) を特定範囲内に制限すると共に、 焼鈍後の冷 却速度を制御することによって、 銅板中の固溶 C， 固溶 Nを調整する方法が、 また 特公昭 61— 45689 号公報には、 Tiど Nbの複合添加によって焼付硬化性を向上させる 方法が開示されている。 For example, in Japanese Patent Application Laid-Open No. 55-141526, Nb is added in accordance with the content of C, N, and A1 in steel, and Nb / (dissolved C + dissolved N) is specified in at%. The method of adjusting the solid solution C and solid solution N in the copper plate by controlling the cooling rate after annealing is described in Japanese Patent Publication No. 61-45689. Improves bake hardenability A method is disclosed.
しかしながら、 上記の鋼板は、 深絞り性に優れる材質とするため、 素材鋼板の強 度は低く、 構造用材料としては必ずしも十分ではない。  However, since the above steel sheet is made of a material having excellent deep drawability, the strength of the material steel sheet is low, and is not always sufficient as a structural material.
また、 特開平 5— 25549 号公報には、 鋼に W, Cr, Moを単独または複合添加する ことによって焼付硬化性を向上させる方法が開示されている。  Also, Japanese Patent Application Laid-Open No. 5-25549 discloses a method of improving bake hardenability by adding W, Cr, Mo alone or in combination to steel.
上記の従来技術において、 焼付硬化により強度が上昇するのは、 鋼板中の微量な 固溶 C , 固溶 Nの働きによるものであり、 また良く知られているように B H鋼板の 場合は材料の降伏強度のみを上昇させるもので、 引張強度を上昇させるものではな い。 従って、 部品の変形開始応力を高める効果しかなく、 変形開始から変形終了 までの変形全域にわたる変形に要する応力 （成形後引張強度） を高める効果は十分 とは言えなかった。  In the above prior art, the increase in strength due to bake hardening is due to the action of trace amounts of solid solution C and solid solution N in the steel sheet, and as is well known, in the case of BH steel sheet, It increases only the yield strength, not the tensile strength. Therefore, the effect of increasing the deformation starting stress of the component was only effective, and the effect of increasing the stress (tensile strength after forming) required for deformation over the entire deformation range from the start of deformation to the end of deformation was not sufficient.
成形後に引張強度が上昇する冷延鋼板として、 例えば特開平 10— 310847号公報に は、 200〜450 の熱処理温度域で引張強度が 60 MPa以上上昇する合金化溶融亜鉛め つき鋼板が開示されている。  As a cold rolled steel sheet whose tensile strength increases after forming, for example, Japanese Patent Application Laid-Open No. 10-310847 discloses an alloyed hot-dip galvanized steel sheet whose tensile strength increases by 60 MPa or more in a heat treatment temperature range of 200 to 450. I have.
この鋼板は、 mass%で、 C ： 0. 01-0. 08% , Mn： 0. 01〜3. 0 %を含有し、 かつ W， Cr, Moの 1種または 2種以上を合計で 0. 05〜3. 0 %含有し、 また必要に応じて Ti : 0. 005 〜0. 1 %, Nb : 0. 005 〜0. 1 %， V： 0. 005 〜0· 1 %の 1種または 2種以上 を含有する組成になり、 かつ鋼のミク口組織がフェライ トまたはフェライト主体か らなるものである。  This steel sheet contains, in mass%, C: 0.01-0.08%, Mn: 0.01-3.0%, and one or more of W, Cr, and Mo in total of 0%. 0.05 to 3.0%, and if necessary, Ti: 0.005 to 0.1%, Nb: 0.005 to 0.1%, V: 0.005 to 0.1% Or, it has a composition containing two or more kinds, and the microstructure of the steel is mainly composed of ferrite or ferrite.
しかしながら、 この技術は、 成形後の熱処理により鋼板中で微細な炭化物を形成 させ、 プレス時に付与する歪みに対して転位を効果的に増殖させて、 歪み量を増加 させるものであるため、 220~370 °Cの温度範囲で熱処理を行う必要があり、 一般 的な焼付硬化処理温度よりも必要とされる熱処理温度が高いという難点があつた。 更に、 昨今の地球環境問題からの排出ガス規制に関連し、 自動車における車体重 量の軽減は極めて重要な課題となっている。 自動車の車体重量軽減のためには、 使 用されている鋼板の強度を増加させ、 すなわち高張力鋼板を適用して、 使用する鋼 板を薄くするのが有効である。 However, this technology forms fine carbides in the steel sheet by heat treatment after forming, effectively increases the dislocation with respect to the strain applied during pressing, and increases the amount of strain. Heat treatment must be performed in the temperature range of 370 ° C, and the required heat treatment temperature is higher than the general bake hardening treatment temperature. Furthermore, in connection with the recent emission control due to global environmental problems, reduction of vehicle weight in automobiles has become a very important issue. In order to reduce the body weight of automobiles, the strength of the steel plate used is increased, that is, the steel used is It is effective to make the board thinner.
薄肉の高張力鋼板を使用した自動車部品は、 その役割に応じた特性が十分に発揮 されねばならない。 特性は、 部品によって異なるが、 例えば耐デント性、 曲げ、 ね じり変形に対する静的強度、 耐疲労性、 耐衝撃特性などがある。 すなわち、 自動車 部品に適用される高張力鋼板は、 成形加工後にかかる特性にも優れることが必要と なる。 これらの特性は、 成形加工後の鋼板の強度に関係するため、 薄肉化を達成す るためには、 使用する高張力鋼板の強度下限を設定する必要がある。  Automotive parts using thin high-strength steel sheets must fully exhibit the characteristics appropriate to their role. Properties vary by part, but include, for example, dent resistance, static strength against bending and torsion, fatigue resistance, and impact resistance. In other words, high-strength steel sheets applied to automotive parts need to have excellent properties after forming. Since these characteristics are related to the strength of the steel sheet after forming, it is necessary to set the lower limit of the strength of the high-strength steel sheet to be used in order to achieve thinning.
一方、 自動車部品を作る過程においては、 鋼板に対してプレス成形が行われる。 プレス成形した場合、 鋼板の強度が高すぎると、 ①形状凍結性が劣化する、 ②延性 が低下し、 成形時に割れやネッキングなどの不具合を生ずる、 また、 板厚を低減し た場合には、 ③耐デント性 （局部的な圧縮荷重負荷により生ずる凹みに対する耐性） が劣化する、 といった問題が生じ、 自動車車体への高張力鋼板の適用拡大を阻んで いた。  On the other hand, in the process of making automobile parts, steel plates are pressed. In the case of press forming, if the strength of the steel sheet is too high, (1) the shape freezing property will deteriorate, (2) the ductility will decrease, and problems such as cracking and necking will occur during forming. Also, if the sheet thickness is reduced, (3) The dent resistance (resistance to dents caused by local compressive load) deteriorates, which has hindered the expansion of the application of high-tensile steel sheets to automobile bodies.
これを打開するための手法として、 例えば外板パネル用の冷延鋼板では、 極低炭 素鋼を素材とし、 最終的に固溶状態で残存する C量を適正範囲に制御した鋼板が知 られている。 この種の鋼板は、 プレス成形時には軟質に保たれ、 形状凍結性、 延性 を確保し、 プレス成形後に行われる、 170 で X 20分程度の塗装焼付工程で起こる歪 時効硬化現象を利用した降伏応力の上昇を得て、 耐デント性を確保しようとするも のである。 この種の鋼板では、 プレス成形時には Cが鋼中に固溶して軟質であり、 —方、 プレス成形後には、 塗装焼付工程で、 プレス成形時に導入された転位に固溶 Cが固着して、 降伏応力が上昇する。  As a method for overcoming this, for example, a cold-rolled steel sheet for outer panel panels is known to use ultra-low carbon steel as a raw material and control the amount of C finally remaining in a solid solution in an appropriate range. ing. This type of steel sheet is kept soft during press forming, secures shape freezing and ductility, and yield stress using the age hardening phenomenon that occurs in the paint baking process at 170 X 20 minutes after press forming is performed after press forming. It seeks to secure dent resistance by gaining an increase. In this type of steel sheet, C forms a solid solution in the steel during press forming and is soft. On the other hand, after press forming, solid solution C adheres to dislocations introduced during press forming in the paint baking process. The yield stress increases.
しかし、 この種の鋼板では、 表面欠陥となるストレーツチヤーストレインの発生 を防止する観点から、 歪時効硬化による降伏応力上昇量は低く抑えられている。 こ のため、 実際に部品の軽量化に寄与するところは小さいという難点があった。  However, in this type of steel sheet, the amount of increase in yield stress due to strain age hardening is kept low from the viewpoint of preventing the occurrence of a strain chain strain that causes surface defects. For this reason, there was a drawback that the part that actually contributed to the weight reduction of components was small.
一方、 外観があまり問題にならない用途に対しては、 固溶 Nを用いて焼付硬化量 をさらに増加させた鋼板や、 組織をフェライ トとマルテンサイ トからなる複合組織 とすることで焼付硬化性をより一層向上させた鋼板が提案されて 、る。 例えば、 特開昭 60-52528号公報には、 C ： 0. 02-0. 15% , Mn： 0. 8 〜3. 5 %、 P ： 0. 02-0. 15% , A1： 0. 10%以下、 N ： 0. 005 〜0. 025 %を含む鋼を 550 。C以下の温 度で卷き取る熱間圧延と、 冷延後の焼鈍を制御冷却熱処理とする延性おょぴスポッ ト溶接性がともに良好な高強度薄鋼板の製造方法が開示されている。 特開昭 60- 5252 8号公報に記載された技術で製造された鋼板は、 フェライ トとマルテンサイ トを主体 とする低温変態生成物相からなる混合組織を有し延性に優れるとともに、 積極的に 添加された Nによる塗装焼付けの際の歪時効を利用して、 高強度を得ようとするも のである。 On the other hand, for applications where the appearance is not a major problem, steel sheets with further increased bake hardening using solid solution N or a composite structure composed of ferrite and martensite Thus, a steel sheet with further improved bake hardenability has been proposed. For example, JP-A-60-52528 discloses that C: 0.02 to 0.15%, Mn: 0.8 to 3.5%, P: 0.02 to 0.15%, A1: 0. 550 steel containing 10% or less, N: 0.005 to 0.025%. A method for producing a high-strength thin steel sheet having both good hot rolling at a temperature of C or lower and controlled cold heat treatment of annealing after cold rolling and good ductility and spot weldability are both disclosed. The steel sheet manufactured by the technique described in JP-A-60-52528 has a mixed structure composed of a low-temperature transformation product phase mainly composed of ferrite and martensite, has excellent ductility, and has a positive effect. The aim is to obtain high strength by using strain aging during baking of paint with added N.
しかしながら、 特開昭 60-52528号公報に記載された技術では、 歪時効硬化による 降伏応力 Y Sの増加量は大きいが引張強さ T Sの増加量が少なく、 また、 降伏応力 Y Sの増加量も大きくばらつくなど機械的性質の変動も大きいため、 現状で要望さ れている自動車部品の軽量化に寄与できるほどの鋼板の薄肉化が期待できない。 また、 特公平 5- 24979 号公報には、 C : 0. 08〜0. 20%、 Mn： 1. 5 〜3. 5 %を含み 残部 Feおよび不可避的不純物からなる成分組成を有し、 組織がフェライ ト量 5 %以 下の均一なべィナイ トもしくは一部マルテンサイ トを含むべィナイトで構成された 焼付硬化性高張力冷延薄鋼板が開示されている。 特公平 5-24979 号公報に記載され た冷延鋼板は、 連続焼鈍後の冷却過程で 400 〜200 °Cの温度範囲を急冷とし、 その 後を徐冷とすることにより、 組織をペイナイ ト主体の組織として、 従来になかった 高い焼付硬化量を得ようとするものである。  However, in the technique described in Japanese Patent Application Laid-Open No. 60-52528, the increase in the yield stress YS due to strain age hardening is large, but the increase in the tensile strength TS is small, and the increase in the yield stress YS is also large. Because of the large fluctuations in mechanical properties, such as variations, it is not possible to expect steel sheets to be thin enough to contribute to the weight reduction of automotive parts currently demanded. In addition, Japanese Patent Publication No. 5-24979 discloses that C: 0.08 to 0.20%, Mn: 1.5 to 3.5%, the composition of which consists of the balance Fe and unavoidable impurities. Discloses a bake-hardenable high-tensile cold-rolled thin steel sheet composed of uniform bainite with a ferrite content of 5% or less or bainite partially containing martensite. In the cold rolled steel sheet described in Japanese Patent Publication No. 5-24979, the structure is mainly made of payite by rapidly cooling the temperature range of 400 to 200 ° C in the cooling process after continuous annealing and then gradually cooling it. The aim is to obtain a higher bake hardening amount than ever before.
しかしながら、 特公平 5- 24979 号公報に記載された鋼板では、 塗装焼付け後に降 伏強さが上昇し、 従来になかった高い焼付け硬化量が得られるものの、 引張強さま では上昇させることができず、 強度部材に適用した場合、 成形後の耐疲労特性、 耐 衝撃特性の向上が期待でき い。 このため、 耐疲労特性、 耐衝撃特性等が強く要求 される用途への適用ができないという問題が残されていた。  However, in the steel sheet described in Japanese Patent Publication No. 5-24979, the yield strength increases after paint baking, and although a higher bake hardening amount than before can be obtained, it cannot be increased with tensile strength. When applied to high-strength members, improvement in fatigue resistance and impact resistance after molding cannot be expected. For this reason, there remains a problem that it cannot be applied to applications that require strong fatigue resistance and impact resistance.
また、 特公昭 61- 12008号公報には、 高い r値を有する高張力鋼板の製造方法が開 示されているが、 この製造方法は、 極低 C鋼を素材として、 冷間圧延後にフェライ ト一オーステナイ ト共存域で焼鈍することに特徴があり、 得られる鋼板が高い r値 と高い塗装焼付け硬化性 （B H性） を有するとされるが、 得られる B H量が高々 60M Pa 程度であり、 また、 この鋼板も時効後に降伏点は上昇するものの、 T Sの上昇は なく、 適用できる部品に限界があるという問題があった。 In addition, Japanese Patent Publication No. 61-12008 discloses a method for producing a high-strength steel sheet having a high r-value. As shown, this production method is characterized by annealing in the coexistence region of ferrite and austenite after cold rolling using ultra-low C steel as a material, and the resulting steel sheet has a high r value and high paint baking. It is said to have hardenability (BH properties), but the amount of BH obtained is at most about 60 MPa. Also, although the yield point of this steel sheet also increases after aging, there is no increase in TS, and the applicable parts are limited. There was a problem that there is.
さらに、 上記した従来の鋼板では、 単純な引張試験による塗装焼付処理後の強度 評価では優れているものの、 実プレス条件にしたがって、 塑性変形させたときの強 度に大きなばらつきが存在し、 信頼性が要求される部品に適用するには必ずしも十 分とはいえなかったのである。  Furthermore, although the above-mentioned conventional steel sheet is excellent in strength evaluation after baking treatment by a simple tensile test, there is a large variation in the strength when plastically deformed according to the actual pressing conditions. It was not necessarily enough to apply to parts that required the above.
プレス成形体の塗装焼付鋼板の中で熱延鋼板に関しては、 例えば特公平 8-23048 号公報に、 加工時には軟質で、 加工後の焼付塗装処理により疲労特性の改善に有効 な引張強度を大幅に上昇させた熱延鋼板の製造方法が開示されている。  Among hot-rolled steel sheets among press-formed baked steel sheets, for example, in Japanese Patent Publication No. 8-23048, the tensile strength, which is soft at the time of processing and is effective in improving fatigue properties by baking coating processing after processing, is greatly increased. A method for producing a raised hot-rolled steel sheet is disclosed.
この技術では、 C量を 0. 02〜0. 13mass%とし、 Nを 0. 0080〜0. 0250mass%と多量 に添加した上で、 仕上圧延温度おょぴ卷取り温度を制御して多量の固溶 Nを鋼中に 残存させ、 金属組織をフェライ トとマルテンサイ トを主体とする複合組織とするこ とで、 成形後熱処理温度： 170 でにて 100 MPa 以上の引張強度の増加が達成される 旨が開示されている。  In this technology, the amount of C is set to 0.02 to 0.13 mass%, N is added to a large amount of 0.0080 to 0.0250 mass%, and then the finish rolling temperature and the winding temperature are controlled to control a large amount. By leaving solid solution N in the steel and forming the metal structure as a composite structure mainly composed of ferrite and martensite, an increase in tensile strength of 100 MPa or more was achieved at a heat treatment temperature of 170 after forming. Is disclosed.
また、 特開平 10— 183301号公報には、 鋼成分のうち、 特に Cと Nを C ： 0. 01〜0. 1 2mass%、 N： 0. 0001〜0. 01mass%に制限すると共に、 平均結晶粒径を 8 111以下に 制御することにより、 80 MPa以上の高 Β Η量を確保すると共に A I量を 45MPa 以下 に抑制することが可能な焼付硬化性およぴ耐室温時効性に優れた熱延鋼板が提示さ れている。  Japanese Patent Application Laid-Open No. 10-183301 discloses that, among the steel components, C and N are particularly limited to C: 0.01 to 0.12 mass%, N: 0.0001 to 0.01 mass%, and By controlling the crystal grain size to 8111 or less, high bake hardenability and high room temperature aging resistance can be achieved, while ensuring a high mass of 80 MPa or more and suppressing the AI amount to 45 MPa or less. Hot rolled steel sheets are presented.
しかしながら、 これらの鋼板は、 熱延板であることから、 仕上圧延後のオーステ ナイ トノフェライ ト変態によりフェライ トの集合組織がランダム化するため、 高 r 値を得ることが困難であり、 十分な深絞り性を有しているとは言い難い。  However, since these steel sheets are hot-rolled sheets, the texture of ferrite is randomized by austenite-noferrite transformation after finish rolling, so that it is difficult to obtain a high r-value, and a sufficient depth is obtained. It is hard to say that it has drawability.
しかも、 これらの技術で得られた熱延鋼板を出発材として冷間圧延および再結晶 焼鈍を行ったとしても、 必ずしも熱延鋼板と同様の成形一熱処理後の引張強度上昇 や 80 MPa以上の高 B Hが得られるとは限らない。 というのは、 鋼組織が、 冷間圧延 およぴ再結晶焼鈍により熱延時とは異なるミクロ組織となること、 また冷間圧延時 に大きな歪蓄積が起こるため、 炭化物、 窒化物または炭窒化物が形成され易く、 固 溶 Cおよぴ固溶 N状態が変化するからである。 In addition, cold rolling and recrystallization using the hot rolled steel sheet obtained by these technologies as starting materials Even if annealed, the same increase in tensile strength after forming and heat treatment as in hot-rolled steel sheets and high BH of 80 MPa or more are not necessarily obtained. This is because the steel structure has a microstructure different from that during hot rolling by cold rolling and recrystallization annealing, and large strain accumulation occurs during cold rolling, so that carbide, nitride or carbonitride Is easily formed, and the state of solid solution C and solid solution N changes.
この発明は、 上記の実状に鑑み開発されたもので、 プレス成形時に優れた深絞り 性を維持しつつ、 プレス成形—熱処理によって引張強度が増加する、 歪時効硬化特 性に優れた冷延鋼板および合金化溶融亜鉛めつき鋼板を、 それらの有利な製造方法 と共に提案することを目的とする。  The present invention has been developed in view of the above-described circumstances, and a cold-rolled steel sheet having excellent strain aging hardening characteristics, in which tensile strength is increased by press forming-heat treatment while maintaining excellent deep drawability during press forming. And to propose alloyed hot-dip galvanized steel sheets together with their advantageous production methods.
また本発明では、 上記従来技術の問題点に鑑み、 T S X r値≥750MPaの優れた深絞 り性おょぴ優れた歪時効硬化特性 （B H≥80MPa かつ A T S≥40MPa ) を有する深 絞り用の冷延鋼板および溶融亜鉛めつき鋼板 （合金化したものも含む） を、 それら の有利な製造方法とともに提供することにある。  Also, in the present invention, in consideration of the above-mentioned problems of the prior art, an excellent deep drawability with a TSX r value ≥750 MPa and an excellent strain age hardening characteristic (BH≥80 MPa and ATS≥40 MPa) are used for deep drawing. It is to provide cold-rolled steel sheets and hot-dip galvanized steel sheets (including alloyed ones) together with their advantageous production methods.
さらに本発明では、 上記した従来技術の問題を解決し、 高度の成形性が要求され る自動車部品用に好適な、 軟質で高い成形性と、 安定した品質特性を有し、 複雑な 形状の自動車部品への成形が容易で、 スプリングパック、 ねじれ、 反り等の形状不 良、 割れ等の発生がないうえ、 さらに自動車部品に成形したのちの熱処理で自動車 部品として十分な強度が得られ自動車車体の軽量匕に充分に寄与できる、 1. 2 以上 という高い r値と、 優れた歪時効硬化特性を有する高張力冷延鋼板およびこれら鋼 板を工業的に安価に、 かつ形状を乱さずに製造できる製造方法を提供することを目 的とする。 発明の開示  Furthermore, the present invention solves the above-mentioned problems of the prior art, and has a soft and high formability and a stable quality characteristic suitable for an automobile part requiring a high degree of formability. It is easy to mold into parts, has no shape defects such as spring pack, torsion, warpage, cracks, etc. In addition, heat treatment after molding into automobile parts provides sufficient strength as automobile parts, and High tensile cold-rolled steel sheets having a high r-value of 1.2 or more and excellent strain aging hardening properties that can sufficiently contribute to lightweight daggers, and these steel sheets can be manufactured industrially at low cost without disturbing the shape. The purpose is to provide a manufacturing method. Disclosure of the invention
本発明者らは、 上記課題を達成するために、 組成おょぴ製造条件を種々変えて鋼 板を製造し、 多くの材質評価実験を行った。 その結果、 高加工性が要求される分野 では従来あまり積極的に利用されることがなかった Nを強化元素として、 この強化 元素の作用により発現する大きな歪時効硬化現象を有利に活用することにより、 成 形性の向上と成形後の高強度化とを容易に両立させることができることを知見した。 さらに、 本発明者らは、 Nによる歪時効硬化現象を有利に活用するためには、 N による歪時効硬化現象を自動車の塗装焼付け条件、 あるいはさらに積極的に成形後 の熱処理条件と有利に結合させる必要があり、 そのために、 熱延条件や冷延、 冷延 焼鈍条件を適正化して、 鋼板の微視組織と固溶 N量とをある範囲に制御することが 有効であることを見いだした。 また、 Nによる歪時効硬化現象を安定して発現させ るためには、 組成の面で、 特に A1含有量を N含有量に応じて制御することが重要で あることも見いだした。 In order to achieve the above object, the present inventors manufactured steel plates with various composition and manufacturing conditions, and performed many material evaluation experiments. As a result, N has been used as a strengthening element in areas where high workability is required. It has been found that by advantageously utilizing the large strain age hardening phenomenon developed by the action of elements, it is possible to easily achieve both improvement in formability and high strength after molding. Furthermore, the present inventors have found that in order to advantageously utilize the strain age hardening phenomenon due to N, the strain age hardening phenomenon due to N is advantageously combined with the baking conditions of automobiles or the heat treatment conditions after molding more positively. For this purpose, it has been found that it is effective to control the microstructure and the amount of solute N in a certain range by optimizing the hot rolling conditions, cold rolling, and cold rolling annealing conditions. . We also found that it is important to control the A1 content according to the N content in terms of composition in order to stably develop the strain age hardening phenomenon due to N.
さらに本発明者らは、 高 r値を得るために、 C含有量を低減し、 フ ライ トーォ ーステナイ トの二相域温度で連続焼鈍を施しその後の冷却を制御して、 フ ライ ト 相中に面積率で 5 %以上のァシキユラ一フェライ ト相を含む組織とし、 このような 微視組織と適正な固溶 N量との組合せで、 高 r値を有しプレス成形性に優れ、 かつ 歪時効硬化特性に優れた冷延鋼板が得られることを見いだした。 また、 これにより、 従来問題であった室温時効劣化の問題もなく、 Nを充分に活用できることを見い出 した。  In order to obtain a high r value, the present inventors further reduced the C content, performed continuous annealing at a temperature in the two-phase region of the fly-to-stenite, and controlled the subsequent cooling to reduce the content in the fly phase. In addition, the microstructure containing 5% or more of the ferrite phase in the area ratio is high, and the combination of such a microstructure and an appropriate amount of dissolved N has a high r value, excellent press formability, and distortion. It has been found that a cold rolled steel sheet having excellent age hardening characteristics can be obtained. In addition, they have found that N can be fully utilized without the problem of room-temperature aging degradation, which was a conventional problem.
すなわち、 本発明者らは、 Nを強化元素として用い、 A1含有量を N含有量に応じ て適正な範囲に制御するとともに、 熱延条件や冷延、 冷延焼鈍条件を適正化して、 微視組織と固溶 Nを最適化することにより、 従来の固溶強化型の C— Mn鋼板、 析出 強化型鋼板に比べて、 高い r値を有し格段に優れた成形性と、 上記した従来の鋼板 にない歪時効硬化特性とを有する鋼板が得られることを見いだしたのである。  That is, the present inventors use N as a strengthening element, control the A1 content in an appropriate range according to the N content, and optimize the hot rolling conditions, cold rolling, and cold rolling annealing conditions, By optimizing the visual structure and solid solution N, the r-value is much higher than the conventional solid solution strengthened C-Mn steel sheet and precipitation hardened It has been found that a steel sheet having strain aging hardening characteristics not found in the above steel sheets can be obtained.
また、 本発明の鋼板は、 単純な引張試験による塗装焼付処理後の強度が従来の鋼 板よりも高いうえ、 さらに実プレス条件にしたがって塑性変形させたときの強度の ばらつきが小さく、 安定した部品強度特性が得られ、 信頼性が要求される部品への 適用が可能となる。 例えば、 歪が大きく加わり板厚が減少した部分は、 他の部分よ り硬化代が大きく （板厚） X (強度） という載荷重能力で評価すると均一化する方 向であり、 部品としての強度は安定するのである。 In addition, the steel sheet of the present invention has a higher strength after a paint baking treatment by a simple tensile test than a conventional steel sheet, and has a small variation in strength when plastically deformed according to actual pressing conditions, and is a stable component. The strength characteristics can be obtained, and it can be applied to parts that require reliability. For example, the part where the plate thickness is reduced due to large strain is larger than other parts, and the hardening allowance is more uniform when evaluated by the load capacity of (plate thickness) X (strength). Direction, and the strength as a part is stable.
発明者らは、 上記の目的を達成すべく更に鋭意研究を重ねた結果、 以下の知見を 得た。  The inventors have further studied diligently to achieve the above object, and have obtained the following findings.
1 ) 成形一熱処理後に引張強度を上昇させるためには、 引張変形を進行させるため に新たな転位を導入する必要がある。 成形により導入された転位と侵入型元素また は析出物との相互作用により、 上降伏応力に達しても予変形により導入された転位 が移動しないことが必要となる。  1) In order to increase tensile strength after forming and heat treatment, it is necessary to introduce new dislocations in order to promote tensile deformation. Due to the interaction between dislocations introduced by forming and interstitial elements or precipitates, it is necessary that the dislocations introduced by pre-deformation do not move even when the upper yield stress is reached.
2 ) W, Cr, Mo, Ti, Nb, Alなどの炭化物、 窒化物または炭窒化物を形成すること によって、 上記の相互作用を得るためには、 成形後の熱処理温度を 200で以上まで 高める必要がある。 従って、 侵入型元素の積極的な活用または Fe炭化物あるいは Fe 窒化物を活用する方が、 成形後の熱処理温度を低下させる点では有利である。  2) To obtain the above interaction by forming carbides, nitrides, or carbonitrides such as W, Cr, Mo, Ti, Nb, and Al, raise the heat treatment temperature after molding to 200 or more to obtain the above interaction. There is a need. Therefore, active utilization of interstitial elements or utilization of Fe carbide or Fe nitride is advantageous in lowering the heat treatment temperature after molding.
3 ) 侵入型元素の中では固溶 Cよりも固溶 Nの方が、 成形後の熱処理温度を低めて も、 成形により導入された転位との相互作用が大きく、 上降伏応力に達しても予変 形に導入された転位が移動し難い。  3) Among the interstitial elements, solid solution N has a larger interaction with dislocations introduced by forming, even if the heat treatment temperature after forming is lowered, even if solid solution N reaches higher yield stress than solid solution C. The dislocations introduced in the predeformation are difficult to move.
4 ) 鋼中の固溶 N存在場所として結晶粒内おょぴ結晶粒界があるが、 成形後の熱処 理以後の強度の増加量は結晶粒界面積が広い方が大きい。 すなわち結晶粒径が小さ い方が有利である。  4) There is a crystal grain boundary in the crystal grain as a place where solid solution N exists in the steel. However, the increase in strength after heat treatment after forming is larger when the crystal grain boundary area is larger. That is, a smaller crystal grain size is advantageous.
5 ) 結晶粒界面積を広くするという観点では、 Nbおよび Bを複合添加すると共に、 熱間圧延終了後直ちに冷却することにより、 熱間圧延終了後のフユライ ト粒の正常 粒成長を抑制し、 かつ冷間圧延に引き続く再結晶焼鈍での粒成長を抑制することが 有利である。  5) From the viewpoint of increasing the grain boundary area, by adding Nb and B in combination and cooling immediately after the completion of hot rolling, normal grain growth of the fluoride grains after the completion of hot rolling is suppressed. It is advantageous to suppress grain growth in recrystallization annealing subsequent to cold rolling.
この発明は、 上記の知見に立脚するものである。 上記の知見は以下の実験から得 られた。 mass%にて、 C ： 0. 0015%, B： 0. 0010%, Si： 0. 01%, Mn： 0. 5 %, P： 0. 03%， S : 0. 008 %および N : 0. 011 %を含み、 かつ Nbを 0. 005〜0. 05%および A1を 0. 0 05-0.03%の範囲で含有し、 残部は Feおよび不可避的不純物の組成になるシートパ - (厚み： 30mm) を、 115(TCで均一加熱した後、仕上温度が Ar3変態点以上の 900。C となるように 3パスで熱間圧延を行い、 圧延終了後、 0.1 秒後に水冷した。 その後、 500で， 1時間のコイル卷取り相当熱処理を実施した。 The present invention is based on the above findings. The above findings were obtained from the following experiments. At mass%, C: 0.0015%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0.008%, and N: 0 0.11%, and Nb 0.005 to 0.05% and A1 0.0 Contained in the range of 05-0.03%, the balance is Fe and unavoidable impurities. The composition of the sheet-pattern (thickness: 30 mm) is 115 (after uniform heating with TC, the finishing temperature is 900 with Ar 3 transformation point or higher). Hot rolling was performed in three passes so as to obtain C. After the rolling was completed, water cooling was performed 0.1 seconds later, and then heat treatment equivalent to coil winding was performed at 500 for 1 hour.
得られた板厚： 4mraの熱延板を、 圧下率： 82.5%で冷間圧延後、 800C, 40秒の 再結晶焼鈍を施し、 ついで圧下率： 0.8 %の調質圧延を施した。 かく して得られ た冷延板から、 圧延方向に JIS 5 号引張試験片を採取し、 通常の引張試験機を用い て、 歪み速度： 0.02/sで引張強度を測定した。 また、 別途、 これらの冷延板から圧 延方向に採取した JIS 5 号引張試験片に 10%の引張歪みを付与し、 120 ， 20分の 熱処理を施したのち、 通常の引張試験に供した。 これら、 冷延板から採取した試験 片の引張強度と 10%の引張り歪を付与後 120°C, 20分の熱処理を行った試験片の引 張強度との差を成形後引張強度上昇代 （ATS) とした。  The obtained hot-rolled sheet having a thickness of 4 mra was cold-rolled at a rolling reduction of 82.5%, then recrystallized and annealed at 800 C for 40 seconds, and then temper-rolled at a rolling reduction of 0.8%. From the cold-rolled sheet thus obtained, a JIS No. 5 tensile test piece was sampled in the rolling direction, and the tensile strength was measured at a strain rate of 0.02 / s using an ordinary tensile tester. Separately, JIS No. 5 tensile test specimens sampled in the rolling direction from these cold-rolled sheets were given a 10% tensile strain, heat-treated for 120 and 20 minutes, and then subjected to a normal tensile test. . The difference between the tensile strength of the test piece taken from the cold-rolled sheet and the tensile strength of the test piece subjected to a heat treatment at 120 ° C for 20 minutes after the application of a 10% tensile strain was determined by the tensile strength increase after forming. ATS).
図 1に、 鋼成分 （N% - 14/93 · Nb% - 14/27 · Al% - 14/11 · B%) と ATSと の関係について調べた結果を示す。  Figure 1 shows the results of a study on the relationship between the steel composition (N%-14/93 · Nb%-14/27 · Al%-14/11 · B%) and ATS.
同図に示したとおり、 （N% - 14/93 - Nb% 一 14/27 · Al% - 14/11 · B%) の 値が 0.0015mass%以上を満足する場合に、 ATSが 60 MPa以上になることが判明した。 実験 2  As shown in the figure, when the value of (N%-14/93-Nb%-14/27 · Al%-14/11 · B%) satisfies 0.0015 mass% or more, the ATS is 60 MPa or more. It turned out to be. Experiment 2
mass%にて、 C： 0.0010%, Si :0.02%， Mn： 0.6 %, P：0.01%, S： 0.009 %， N：0.012 %, A1:0.01%および Nb: 0.015 %を含み、かつ Bを 0.00005 〜0.0025% の範囲で含有し、 残部は Feおよび不可避的不純物の組成になるシートパー （厚み： 3 0mm) を、 1100°Cで均一加熱したのち、 仕上温度が Ars変態点以上の 920 ^となるよ うに 3パス圧延を行い、 圧延終了後、 0.1 秒後に水冷し、 コイル卷取り相当熱処理 を 450°C， 1時間実施した。  In mass%, C: 0.0010%, Si: 0.02%, Mn: 0.6%, P: 0.01%, S: 0.009%, N: 0.012%, A1: 0.01% and Nb: 0.015%, and B The content is in the range of 0.00005 to 0.0025%, and the remainder is a sheet par (thickness: 30 mm) with the composition of Fe and unavoidable impurities. After heating uniformly at 1100 ° C, the finishing temperature is 920 ^ above the Ars transformation point. Three-pass rolling was performed as much as possible, and water cooling was performed 0.1 seconds after the rolling was completed, and heat treatment equivalent to coil winding was performed at 450 ° C for 1 hour.
得られた板厚： 4ramの熱延板を、 圧下率： 82.5%で冷間圧延後、 820。C, 40秒の再 結晶焼鈍を施し、 ついで圧下率： 0.8 %の調質圧延を施した。  The obtained hot-rolled sheet having a thickness of 4 ram was cold-rolled at a reduction ratio of 82.5%, and then 820. C, recrystallization annealing for 40 seconds, followed by temper rolling at a reduction of 0.8%.
かくして得られた冷延板から、 圧延方向に JIS 5号引張試験片を採取し、 通常の 引張試験機を用いて、 歪み速度： 0. 02/sで引張強度を測定した。 また、 別途、 これ らの冷延板から採取した引張試験片に 10%の引張歪みを付与し、 120 ， 20分の熱処 理を施したのち、 通常の引張試験に供した。 From the cold-rolled sheet thus obtained, take a JIS No. 5 tensile test specimen in the rolling direction and The tensile strength was measured at a strain rate of 0.02 / s using a tensile tester. Separately, 10% tensile strain was applied to tensile test specimens collected from these cold-rolled sheets, heat-treated for 120 and 20 minutes, and then subjected to normal tensile tests.
図 2に、 鋼中の B含有量と A TSとの関係について調べた結果を示す。 同図に示 したとおり、 Bを 0. 0005〜0. 0015mass%含有する場合に 60 MPa以上の高い Δ TSが得 られることが分かる。  Figure 2 shows the results of a study on the relationship between the B content in steel and ATS. As shown in the figure, when B is contained in 0.0005 to 0.0015 mass%, a high ΔTS of 60 MPa or more can be obtained.
また、 Nbと Bを複合添加することによって結晶粒が微細化され、 高い A TSが得ら れることがミク口組織観察により判明した。  In addition, it was found from microstructure observation that the addition of Nb and B in combination resulted in the refinement of the crystal grains and high ATS.
すなわち、 B量が 0. 0005mass%未満では Nbとの複合添加による結晶粒微細化効果 が小さい。 逆に B量が 0. 0015mass%を超える場合には、 粒界おょぴその近傍に偏析 する B量が増加し、 かかる B原子は N原子間との相互作用が強いことから有効な固 溶 N量が低下するため ATSが低下したものと推察される。  That is, if the B content is less than 0.0005 mass%, the effect of refining crystal grains by adding Nb in combination is small. Conversely, if the B content exceeds 0.0015 mass%, the amount of B segregating in the vicinity of the grain boundary increases, and such B atoms have a strong interaction with N atoms, so that effective B It is probable that ATS decreased due to the decrease in N content.
実験 3 Experiment 3
mass%にて、 C ： 0. 0010%, N ： 0. 012 %, B ： 0. 0010%, Si： 0. 01%, Mn： 0. 5 %, P : 0. 03%, S : 0. 008 %, Nb： 0. 014 %および A1 : 0. 01%を含有し、 残部は Feおよび不可避的不純物の組成になる鋼 Aと、 C : 0. 01 0 %, N ： 0. 0012%, B ： 0. 0010%, Si ： 0. 01 %, Mn： 0. 5 %, P ： 0. 03%, S ： 0, 008 %, Nb： 0, 014 %お ょぴ A1 ： 0. 01%を含有し、 残部は Feおよび不可避的不純物の組成になる鋼 Bの各シ ートパー （厚み： 30ram) を、 1150 で均一加熱した後、 仕上温度が Ar3変態点以上 の 910でとなるように 3パス圧延を行い、 圧延終了後、 0. 1秒後にガス冷却を開始し、 引き続き 600 で 1時間のコイル卷取り相当熱処理を実施した。 At mass%, C: 0.0010%, N: 0.012%, B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0 008%, Nb: 0.014% and A1: 0.011%, the balance being Fe and the composition of inevitable impurities A, C: 0.010%, N: 0.0012% , B: 0.0010%, Si: 0.01%, Mn: 0.5%, P: 0.03%, S: 0, 008%, Nb: 0, 014% A1: 0.01 %, The balance being Fe and unavoidable impurities. After uniformly heating each sheeter (thickness: 30 ram) of steel B at 1150, the finishing temperature will be 910 which is above the Ar 3 transformation point. After the rolling was completed, gas cooling was started 0.1 seconds after the rolling was completed, and heat treatment equivalent to coil winding was performed at 600 ° C for 1 hour.
得られた板厚： 4 nunの熱延板を、 圧下率 82. 5%で冷間圧延したのち、 880¾, 40 秒の再結晶焼鈍を施し、 ついで圧下率： 0. 8 %の調質圧延を施した。  The resulting hot-rolled sheet with a thickness of 4 nun is cold-rolled at a rolling reduction of 82.5%, then recrystallized and annealed at 880¾ for 40 seconds, and then temper-rolled with a rolling reduction of 0.8%. Was given.
かく して得られた冷延板から、 圧延方向に JIS 5号引張試験片を採取し、 通常の 引張試験機を用いて、 歪み速度： 0. 02/sで引張強度を測定した。 また、 別途、 これ らの冷延板から採取した引張試験片に 10%の引張歪みを付与し、 種々の温度で 20分 間の熱処理を施したのち、 通常の引張試験に供した。 From the cold-rolled sheet thus obtained, a JIS No. 5 tensile test specimen was sampled in the rolling direction, and the tensile strength was measured at a strain rate of 0.02 / s using an ordinary tensile tester. Separately, a 10% tensile strain was applied to tensile test specimens collected from these cold-rolled sheets, and the specimens were treated at various temperatures for 20 minutes. After the heat treatment, the steel sheet was subjected to a normal tensile test.
図 3に、 に及ぼす成形後熱処理温度の影響について調べた結果を示す。 同 図に示したとおり、 成形後熱処理温度が 200 以下と比較的低い領域では極低炭、 高 N含有鋼である鋼 Aの方が、 セミ極低炭 ·低 N鋼である鋼 Bよりも高い ATSを示 し、 高温域では同程度の ATSを示す。 これらの実験結果から、 低温域での ATSを 確保するには固溶 Nを活用すること  Figure 3 shows the results of an investigation on the effect of the post-molding heat treatment temperature on. As shown in the figure, in a region where the post-forming heat treatment temperature is relatively low at 200 or less, steel A, which is a steel with extremely low carbon and high N content, is more than steel B, which is a semi-ultra low carbon and low N steel. It shows high ATS, and shows similar ATS at high temperature. From these experimental results, it is necessary to use solid solution N to secure ATS at low temperatures.
が有効であることが分かる。 Is effective.
また、 図 4に、 常温時効による伸びの低下量 （ΔΕ1) と成形後引張強度上昇代 （Δ TS) に及ぼす、結晶粒径 dと鋼成分 （N% - 14/93 - Nb% 一 14/27 - Al% - 14/11. B%) との影響について調べた結果を示す。 なお、 伸びの低下量 （ΔΕ1) は、 冷延 板から圧延方向に採取した JIS 5 号試験片で測定した全伸びと、 別途採取した試験 片を用い常温時効の促進処理である 100でで 8時間の保持処理を施したのちに測定 した全伸びとの差で評価した。  Fig. 4 shows the effect of crystal grain size d and steel composition (N%-14/93-Nb%-14 /) on the reduction in elongation due to normal temperature aging (Δ 代 1) and the increase in tensile strength after forming (ΔTS). 27-Al%-14/11. B%). The amount of decrease in elongation (ΔΕ1) was calculated using the total elongation measured with a JIS No. 5 test piece taken in the rolling direction from the cold-rolled sheet and 100, which is a normal temperature aging acceleration treatment using a separately taken test piece. The evaluation was made based on the difference from the total elongation measured after the time holding treatment.
同図に示したとおり、 （N% - 14/93 - Nb% - 14/27 - Al% - 14/11 · B%) の 値が 0.0015mass%以上でかつ結晶粒径 dが 20/ m 以下の場合に、 高 ATSと低 ΔΕ1の 両立が可能となることが分かる。  As shown in the figure, the value of (N%-14/93-Nb%-14/27-Al%-14 / 11B%) is 0.0015 mass% or more and the crystal grain size d is 20 / m or less. It can be seen that in the case of, both high ATS and low ΔΕ1 can be achieved.
実験 4 Experiment 4
0.0015%C-0.30%Si-0.8%Mn-0.03 P~0.005%S-0.012%N- 0.02〜0· 08A1鋼のシートパ 一を 1050 に均一に加熱し、 次いで仕上温度が 670。Cになるように 7パスで熱間仕 上圧延し、 次いで 700。CX 5h の再結晶焼鈍を行い、 得られた板厚 4mmの熱延板を、 圧下率 82.5%で冷間圧延し、 次いで 875 で X 40秒で再結晶焼鈍し、 次いで圧下率 0.8 %で調質圧延し、 得られた冷延板から JIS 5号引張試験片を採取し、 通常の引張試 験機を用いて歪速度 3 X10— 3/sで引張試験を行い、 TS X r値および ATSを測定 した。 結果を表 5に示す。 Ν/Α1≥0· 30を満足する場合に T S X r値≥750 かつ Δ Τ S≥40MPa が達成されている。 なお、 NZA1≥0.30の場合、 BH≥80MPa が達成さ れていることは別途確認した。 実験 5 0.0015% C-0.30% Si-0.8% Mn-0.03 P ~ 0.005% S-0.012% N- 0.02-0.08 A1 steel sheet is uniformly heated to 1050, then finishing temperature is 670. Hot finish rolling in 7 passes to C, then 700. CX 5h was recrystallized and annealed, and the obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled at a rolling reduction of 82.5%, then recrystallized at 875 X for 40 seconds, and then adjusted at a rolling reduction of 0.8%. and quality rolled, resulting JIS 5 No. tensile test pieces were taken from cold-rolled sheet, subjected to tensile tests at a strain rate of 3 X10- 3 / s using a conventional tensile test machine, TS X r values and ATS Was measured. Table 5 shows the results. When Ν / Α1≥0 · 30 is satisfied, TSX r value ≥750 and Δ ΤS≥40MPa are achieved. When NZA1≥0.30, it was separately confirmed that BH≥80MPa was achieved. Experiment 5
\  \
0.0015½C-0.0010°/oB-0.01%Si-0.5% n-0.03 P-0.008 S-0.011%N- 0.005 〜0.05%Nb- 0.005 〜0.03%A1 鋼のシートパーを 1000°Cに均一に加熱し、 次いで仕上温度が 650 °Cになるように 7パスで熱間仕上圧延し、 次いで 800。CX60秒の再結晶焼鈍を行い、 得られた板厚 4 mmの熱延板を、 圧下率 82.5%で冷間圧延し、 次いで 880 °CX 40秒で再結晶焼鈍し、 次いで圧下率 0.8 %で調質圧延し、 得られた冷延板から：！  0.0015 ° C-0.0010 ° / oB-0.01% Si-0.5% n-0.03 P-0.008 S-0.011% N- 0.005 to 0.05% Nb- 0.005 to 0.03% A1 Heat the steel sheet par uniformly to 1000 ° C, Next, hot finish rolling is performed in 7 passes so that the finishing temperature becomes 650 ° C, and then 800. Recrystallization annealing was performed for 60 seconds at CX, and the obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled at a reduction of 82.5%, then recrystallized at 880 ° C for 40 seconds, and then reduced at a reduction of 0.8%. From the cold rolled sheet obtained by temper rolling and:
IS 5号引張試験片を採取し、通常の引張試験機を用いて歪速度 3 ΧΙΟ—3/sで引張試 験を行い、 TS X r値、 BH、 ATSを測定した。 これら測定値と N/ (Al+Nb+B) と の関係を図 5に示す。 本実験では Nb： 0.005 -0.05%, B ： 0.0010%を含有してい る鋼を用いており、 同図に示すように、 N/ (Al+Nb+B) ≥0.30の範囲で B H≥80MPa 、 厶 TS≥60MPa 、 T S X r値≥850MPaが達成された。 IS-5 tensile test specimens were collected and subjected to a tensile test using an ordinary tensile tester at a strain rate of 3ΧΙΟ-3 / s, and the TS Xr value, BH, and ATS were measured. Figure 5 shows the relationship between these measured values and N / (Al + Nb + B). In this experiment, steel containing Nb: 0.005 -0.05% and B: 0.0010% was used. As shown in the figure, BH≥80MPa, N / (Al + Nb + B) ≥0.30, TS≥60MPa and TSX r value≥850MPa were achieved.
実験 6 Experiment 6
0.0010%C-0.02%Si-0.6%Mn-0.01%P-0.009%S-0.015%N-0.01%A1— 0.015%Nb-0.0001〜 0.0025%B鋼のシートパーを 1050 に均一に加熱し、 次いで仕上温度が 680 。Cになる ように 7パスで熱間仕上圧延し、 次いで 750 °CX 5h のパッチ焼鈍にて再結晶焼鈍 を行い、 得られた板厚 4 mmの熱延板を、 圧下率 82.5%で冷間圧延し、 次いで 880 °C X40秒で再結晶焼鈍し、 次いで圧下率 0.8 %で調質圧延し、 得られた冷延板から JIS 0.0010% C-0.02% Si-0.6% Mn-0.01% P-0.009% S-0.015% N-0.01% A1-0.015% Nb-0.0001 ~ 0.0025% B Heat the steel sheet par uniformly to 1050, then finish The temperature is 680. C, hot finish rolling in 7 passes, followed by recrystallization annealing by patch annealing at 750 ° C for 5 h.The resulting hot-rolled sheet with a thickness of 4 mm was cold-rolled at a rolling reduction of 82.5%. Rolled, then recrystallized and annealed at 880 ° C for 40 seconds, and then temper rolled at a reduction of 0.8%.
5号引張試験片を採取し、 通常の引張試験機を用いて歪速度 3 X10— /sで引張試 験を行い、 TSX r値、 BH、 AT Sを測定した。 これらの測定値と B量との関係 を図 6に示す。 同図に示すように、 B ： 0.0003~0.0015%の範囲で BH≥80MPa に加え、 Bく 0.0 003%の場合よりも高い Δ T Sレベルである Δ T S≥60MPa、 T S X r値≥850MPaが 達成された。 また、 ミクロ組織観察から、 この B量範囲で結晶粒がとくに微細化し ていることが認められた。 実験 5、 6の結果から、 NZ (Al + Nb+B) ≥0.30の範囲として B≥0, 0003%と し、 さらに Nbの複合添加により結晶粒が微細化し、 ATS、 TS X r値レベルがさ らに改善されることが判明した。 B < 0. 0003%では Nbとの複合添加による結晶粒微 細化効果がない。 一方、 B〉0. 0015%ではかえって特性が低下する。 これは、 粒界 およぴその近傍に偏析する B量が増加し、 B原子と N原子間の強い相互作用により 有効な固溶 N量が低減したためと推察される。 また、 Nbの代わりに Ti、 Vを添加し た場合についても同様の検討を行い、 Nbと同様の効果が得られることを確認した。 本努明は、 以上の知見に基づいてなされたものであり、 その要旨は以下の通りである。 第 1の本発明は mass %で、 A No. 5 tensile test piece was sampled and subjected to a tensile test using an ordinary tensile tester at a strain rate of 3 × 10— / s, and the TSX r value, BH, and ATS were measured. Figure 6 shows the relationship between these measured values and the B content. As shown in the figure, B: In addition to BH≥80MPa in the range of 0.0003% to 0.0015%, ΔTS levels of ΔTS≥60MPa and TSX r value≥850MPa, which are higher than those of B and 0.0003%, were achieved. Was. From microstructural observation, it was confirmed that crystal grains were particularly fine in this B content range. From the results of Experiments 5 and 6, it was assumed that B≥0, 0003% in the range of NZ (Al + Nb + B) ≥0.30. Sa It was found to be improved. When B <0.0003%, there is no crystal grain refinement effect due to the complex addition with Nb. On the other hand, when B> 0.0015%, the characteristics are rather deteriorated. This is presumably because the amount of B segregating at and near the grain boundary increased, and the effective amount of solute N decreased due to the strong interaction between the B and N atoms. In addition, the same examination was performed for the case where Ti and V were added instead of Nb, and it was confirmed that the same effect as Nb was obtained. This effort was made based on the above findings, and the summary is as follows. The first invention is mass%,
C ： 0. 15 %以下、 C: 0.15% or less,
Si ： 1. 0 %以下、 Si: 1.0% or less,
Mn ： 2. 0 %以下、 Mn: 2.0% or less,
P ： 0. 1%以下、 P: 0.1% or less,
S ： 0. 01%以下 S: 0.01% or less
A1 ： 0. 005 〜0. 030 %、 A1: 0.005 to 0.030%,
N ： 0. 0050 〜0. 0400%、 N: 0.0050 to 0.0400%,
を含み、 かつ N/A1： 0. 30以上、 And N / A1: 0.30 or more,
固溶状態の Nが 0. 0010%以上あり、 0.0010% or more of N in solid solution state,
残部が Feおよび不可避的不純物からなる組成を有することを特徴とする歪時効硬化 特性に優れた冷延鋼板。 A cold-rolled steel sheet having excellent strain aging hardening characteristics, characterized in that the balance is composed of Fe and unavoidable impurities.
第 1の本発明では前記組成の中でも特に下記の範囲が好ましい。 すなわち mass%で C ： 0. 01%未満、  In the first present invention, the following ranges are particularly preferable among the above compositions. That is, C: less than 0.01% in mass%,
Si： 0. 005 ~1. 0 %、 Si: 0.005 to 1.0%,
Mn： 0. 01— 1. 5%, Mn: 0.01-1.5%,
P ： 0. 1%以下， P: 0.1% or less,
S ： 0. 01%以下 S: 0.01% or less
A1： 0. 005 ~0. 030 %、 A1: 0.005 to 0.0030%,
N： 0. 005 〜0. 040 %、 を含み、 かつ N/Al： 0.30以上、 N: 0.005 to 0.440%, And N / Al: 0.30 or more,
固溶状態の Nが 0.0010%以上あり、 0.0010% or more of N in solid solution state,
残部が Feおよび不可避的不純物からなる組成を有することを特徴とする歪時効硬化 特性に優れた冷延鋼板。 A cold-rolled steel sheet having excellent strain aging hardening characteristics, characterized in that the balance is composed of Fe and unavoidable impurities.
第 1の本発明では、 上記組成に加えてさらに、 mass%で  In the first present invention, in addition to the above composition,
B ： 0.0001〜0.0030%、 B: 0.0001-0.0030%,
Nb： 0.005 〜0.050 %、 Nb: 0.005 to 0.050%,
を、 次式（1), (2) With the following equations (1), (2)
N%≥0.0015 + 14/93 · Nb% + 14/27 - Al% + 14/11 · B% —- (1) C%≤ 0.5 · (12/93) - Nb% --- (2) を満足する範囲で含むことが好ましい。  N% ≥0.0015 + 14 / 93Nb% + 14/27-Al% + 14 / 11B% --- (1) C% ≤ 0.5 (12/93)-Nb% --- (2) It is preferable to include it in a satisfying range.
第 1の本発明では、 上記組成に加えてさらに、 mass%で必要に応じて Cu、 Ni、 Mo のうちの 1種または 2種以上を合計で 1.0 %以下含むことが好ましい。  In the first aspect of the present invention, it is preferable that, in addition to the above composition, one or more of Cu, Ni, and Mo be contained in a mass% of 1.0% or less as necessary.
第 1の本発明では、 鋼板の結晶粒径が 20 /zm以下であることが好ましい。  In the first aspect of the present invention, the steel sheet preferably has a crystal grain size of 20 / zm or less.
第 1の本発明では、 熱処理温度： 120 〜200 での低温域にて、 成形後の強度上昇 代： 60 MPa以上を有することが好ましい。  In the first aspect of the present invention, it is preferable that the strength increase after molding is 60 MPa or more in the low temperature range of heat treatment temperature: 120 to 200.
第 1の本発明では、 上記冷延鋼板の表面に、 電気亜鉛めつき、 溶融亜鉛めつき、 および合金化溶融亜鉛めつき層をそなえても良い。 第 2の本発明は、 mass%にて  In the first aspect of the present invention, the surface of the cold-rolled steel sheet may be provided with an electrogalvanized layer, a hot-dip galvanized layer, and an alloyed hot-dip galvanized layer. The second invention is based on mass%
C： 0.01%未満、  C: less than 0.01%,
Si： 0.005 ~1.0%、  Si: 0.005 to 1.0%,
Mn： 0.01〜1.5%、  Mn: 0.01-1.5%,
P ： 0.1%以下、  P: 0.1% or less,
S ： 0.01%以下  S: 0.01% or less
A1： 0.005 〜0.030%、 N： 0. 005 〜0. 040%、 0 A1: 0.005 to 0.030%, N: 0.005 to 0.404%, 0
を含み、 かつ N/Al： 0. 30以上 And N / Al: 0.30 or more
を満足する範囲において含有し、 残部は実質的に Feの組成になる鋼片を、 熱間圧延 し、 その際、 仕上圧延終了後直ちに冷却を開始して卷取り温度： 400 〜800 。Cで卷 取り、 その後圧下率： 60〜95%の冷間圧延を施したのち、 650〜900 の温度で再 結晶焼鈍を施すことを特徴とする、 歪み時効硬化特性に優れた冷延鋼板の製造方法 である。 The steel slab having a composition of substantially Fe is hot-rolled. At that time, cooling is started immediately after finishing rolling, and the winding temperature is 400 to 800. C rolled, then cold-rolled at a reduction rate of 60-95%, and then re-crystallized at 650-900 to produce a cold-rolled steel sheet with excellent strain age hardening characteristics. Manufacturing method.
第 2の本発明では、 上記組成に加えてさらに、 mass%で  In the second invention, in addition to the above composition,
B ： 0. 0001~0. 0030%、 B: 0.0001 to 0.0030%,
Nb： 0. 005 〜0. 050 %、 Nb: 0.005 to 0.050%,
を、 次式（1) , (2) With the following equations (1) and (2)
N %≥0. 0015 + 14/93 - Nb% + 14/27 - Al% + 14/11 · B % --- (1) C %≤ 0. 5 - (12/93) - Nb% —- (2) を満足する範囲で含むことが好ましい。  N% ≥0.0015 + 14/93-Nb% + 14/27-Al% + 14 / 11B% --- (1) C% ≤ 0.5-(12/93)-Nb% --- It is preferable to include (2) in a range that satisfies (2).
第 2の本発明では、 上記した再結晶焼鈍における昇温過程において、 500 から 再結晶温度までの温度域を 1〜20eC/sの速度で昇温することが好ましい。 In the second invention, in the Atsushi Nobori process in the recrystallization annealing as described above, it is preferable to raise the temperature of the temperature range up to the recrystallization temperature to 500 at a rate of 1~20 e C / s.
第 2の本発明では、 再結晶焼鈍後、 溶融亜鉛めつき処理、 ついで加熱合金化処理 を施してもよい。 第 3の本発明は、 mass%で、  In the second aspect of the present invention, after the recrystallization annealing, a hot-dip galvanizing treatment and then a heat alloying treatment may be performed. The third aspect of the present invention provides a
C ： 0. 01%以下、  C: 0.01% or less,
Si ： 1. 0 %以下、 Si: 1.0% or less,
n ： 0. 01〜； 1. 5 %、  n: 0.01 to; 1.5%,
P ： 0. 1%以下、  P: 0.1% or less,
S ： 0. 01%以下  S: 0.01% or less
A1 ： 0. 005 〜0. 020 %、 N ： 0.0050 〜0.040%、 A1: 0.005 to 0.002%, N: 0.0050 to 0.040%,
を含み、 かつ N/Al： 0.30以上、 固溶 N： 0.0010%以上で残部 Feおよぴ不可避的不純 物からなる組成を有し、 N / Al: 0.30 or more, solid solution N: 0.0010% or more, with the balance being Fe and unavoidable impurities,
TS X r値： 750MPa以上であることを特徴とする歪時効硬化特性に優れた深絞り用 冷延鋼板である。  TS Xr value: Cold-rolled steel sheet for deep drawing with excellent strain aging hardening characteristics characterized by being 750 MPa or more.
第 3の本発明では、 上記組成に加えてさらに、 mass%で  In the third invention, in addition to the above composition,
B ： 0·0001〜0·0030%、 B: 00001 to 00030%,
Nb： 0.005 〜0.Q50 %、 Nb: 0.005-0.Q50%,
を、 次式（1)， (2) With the following equations (1) and (2)
N%≥0.0015 + 14/93 · Nb% + 14/27 ' Al% + 14/11 · B% --- (1) C%≤ 0.5 · (12/93) - Nb% -— (2) を満足する範囲で含むことが好ましい。'  N% ≥0.0015 + 14 / 93Nb% + 14/27 'Al% + 14 / 11B% --- (1) C% ≤ 0.5 (12/93)-Nb% ---- (2) It is preferable to include it in a satisfying range. '
第 3の本発明では、 上記組成に加えてさらに、 raass%で  In the third invention, in addition to the above composition,
B ： 0.0001〜0.0030%、  B: 0.0001-0.0030%,
Nb： 0.005 —0.050 %、  Nb: 0.005-0.050%,
Ti： 0.005〜0.070 %、  Ti: 0.005-0.070%,
V： 0.005 〜0.10%,  V: 0.005 to 0.10%,
のうち 1種または 2種以上を含有し、 Containing one or more of
かつ N/(A1+Nb+Ti+V+B) : 0.30以上、 And N / (A1 + Nb + Ti + V + B): 0.30 or more,
固溶 N： 0.0010%以上 Solid solution N: 0.0010% or more
を満足する範囲で含むことが好ましい。 第 4の本発明は、 mass%で、 ' Is preferably contained in a range satisfying the following. The fourth invention is based on mass%
C ： 0.01 %以下、 C: 0.01% or less,
Si： 0.005〜： 1.0%、Si: 0.005 ~: 1.0%,
n： 0.01〜1.0 %、 P ： 0. 1%以下、 n: 0.01-1.0%, P: 0.1% or less,
S ： 0. 01%以下  S: 0.01% or less
Al： 0. 005 〜0. 030 %、  Al: 0.005 to 0.030%,
N： 0. 0050 〜0. 040%、  N: 0.0050-0.040%,
を含み、 Including
B ： 0. 0003〜0. 0030%、  B: 0.0003-0.0030%,
Nb： 0. 005 〜0· 050%、  Nb: 0.005 to 0.050%,
Ti： 0. 005 〜0. 070 %、  Ti: 0.005 to 0.70%,
V： 0. 005 〜0. 10%、  V: 0.005 to 0.10%,
のうち 1種または 2種以上を含み、 かつ One or more of the above, and
N/ (Al+Nb+Ti+V+B) ： 0. 30以上になる組成を有する鋼素材を、 950 t以上に加熱後、 粗圧延終了温度を 1000°C以下 Ar3以上として粗圧延し続い Ar3以下 600 °C以上の温 度域で潤滑しつつ仕上圧延して巻き取り、 その際粗圧延開始から仕上圧延終了まで の全圧下率を 80%以上とし、得られた熱延板を再結晶焼鈍し、次いで圧下率 60〜95% で冷間圧延し、 得られた冷延板を再結晶焼鈍することを特徴とする歪時効硬化特性 に優れた深絞り用冷延鋼板の製造方法である。 第 5の本発明 fま、 mass%で、 N / (Al + Nb + Ti + V + B): A steel material having a composition of 0.30 or more is heated to 950 t or more, and then rough-rolled at an end temperature of 1000 ° C or less and Ar 3 or more. Subsequently, finish rolling and winding are performed while lubricating in the temperature range of Ar 3 or less and 600 ° C or more.At that time, the total rolling reduction from the start of rough rolling to the end of finish rolling is set to 80% or more, and the obtained hot rolled sheet A method for producing a cold-rolled steel sheet for deep drawing with excellent strain aging hardening characteristics, characterized by recrystallization annealing and then cold rolling at a rolling reduction of 60 to 95%, and recrystallization annealing the obtained cold-rolled sheet. It is. Fifth invention f
C ： 0. 0015〜0. 025 %、 C: 0.0015 to 0.025%,
Si： 1. 0 %以下、 Si: 1.0% or less,
Mn： 2. 0 %以下、 Mn: 2.0% or less,
P ： 0. 1%以下、 P: 0.1% or less,
S ： 0. 02%以下、 S: 0.02% or less,
A1： 0. 02%以下、 A1: 0.02% or less,
N： 0. 0050〜0. 0250%、 N: 0.0050-0.0250%,
を含み、 かつ B： 0.0005〜0· 0050%、 Including, and B: 0.0005-0.0050%,
Nb： 0.002〜0.050 %、 Nb: 0.002 to 0.050%,
の 1種または 2種以上を含み、かつ N/A1が 0.3 以上、固溶状態としての Nを 0.0010% 以上含有し、 残部が Feおよび不可避的不純物からなる組成と、 面積率で 5%以上の ァシキユラ一フェライ ト相と平均結晶粒径： 20/xm 以下のフェライト相から成る組 織を有し、 r値： 1.2 以上であることを特徴とする成形性、 歪時効硬化特性および 耐常温時効性に優れた高張力冷延鋼板である。 At least 0.3% of N / A1 and 0.0010% or more of N as a solid solution, with the balance being Fe and unavoidable impurities. Formability, strain aging hardening characteristics and room temperature aging resistance characterized by having an organization consisting of a ferrite phase and a ferrite phase having an average crystal grain size of 20 / xm or less and an r value of 1.2 or more. It is a high-tensile cold-rolled steel sheet with excellent resistance
第 5の本発明では、 前記組成に加えてさらに、 mass%で、 下記 a群〜 c群のうち の 1群または 2群以上を含むことが好ましい。  In the fifth aspect of the present invention, it is preferable that one or more of the following groups a to c be further contained in mass% in addition to the composition.
記  Record
a群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1.0 %以下 b群：. Ti、 Vのうちの 1種または 2種を合計で 0.1 %以下  Group a: 1.0% or less in total of one or more of Cu, Ni, Cr, and Mo Group b: 0.1% or less in total of one or two of Ti and V
c群： Ca、 RE の 1種または 2種を合計で 0.0010〜0.010 % 第 6の本癸明は、 raass%で、  Group c: One or two of Ca and RE are 0.0010 to 0.010% in total.
C： 0.0015〜0.025。 C: 0.0015 to 0.025.
Si： 1.0 %以下、Si: 1.0% or less,
n： 2.0 %以下、  n: 2.0% or less,
P： 0.1%以下、 P: 0.1% or less,
S ： 0.02%以下、 S: 0.02% or less,
A1： 0.02%以下、 A1: 0.02% or less,
N： 0.0050〜0.0250% N: 0.0050-0.0250%
B： 0.0003〜0.0050%、 B: 0.0003-0.0050%,
Nb： 0.002 —0.050 %、 Nb: 0.002-0.050%,
の 1種または 2種以上を含み、 かつ N/A1が 0.3 以上である組成の鋼スラブを、 スラブ加熱温度： 1000°C以上に加熱し、 粗圧延してシートパーとなし、 A steel slab containing one or more of the following and having a N / A1 of 0.3 or more is heated to a slab heating temperature of 1000 ° C or more, Rough rolling and sheet par, none
該シートバーに仕上圧延出側温度： 800 以上とする仕上圧延を施し、 Subjecting the sheet bar to finish rolling at a finish rolling exit temperature of 800 or more;
卷取温度： 650 °C以下で卷き取り熱延板とする熱間圧延工程と、 Winding temperature: hot rolling process at 650 ° C or lower to make hot rolled sheet;
該熱延板に酸洗および冷間圧延を施し冷延板とする冷間圧延工程と、 A cold rolling step in which the hot-rolled sheet is subjected to pickling and cold rolling to form a cold-rolled sheet,
該冷延板にフェライトーオーステナイ トニ相域内の温度で連続焼鈍を行い、Continuous annealing is performed on the cold-rolled sheet at a temperature within the ferrite-austenite-tonite region,
500 ¾以下の温度域まで冷却速度： 10〜 300 / sで冷却する冷延板焼鈍工程とを、 順次施すことを特徴とする The cooling rate is 10 to 300 / s to a temperature range of 500 ° C or less.
r値： 1. 2 以上を有し、 成形性、 歪時効硬化特性おょぴ耐常温時効性に優れた高張 カ冷延鋼板の製造方法である。 r-value: This is a method for producing cold-rolled steel sheets with excellent formability, strain aging hardening properties, and normal-temperature aging resistance, having at least 1.2.
第 6の本発明では、 前記組成に加えてさらに、 mass%で、 下記 a群〜 c群のうち の 1群または 2群以上を含むことが好ましい。  In the sixth aspect of the present invention, it is preferable that one or more of the following groups a to c be further contained in mass% in addition to the composition.
記  Record
a群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0 %以下、 b群： Ti、 Vのうちの 1種または 2種以上を合計で 0. 1 %以下 Group a: 1.0% or less in total of one or more of Cu, Ni, Cr, Mo, Group b: 0.1% in total of one or more of Ti, V Less than
c群： Ca、 REM の 1種または 2種を合計で 0. 0010~0. 010 % 第 7の本癸明は、 raass%で、 Group c: Ca, REM 1 or 2 types in total 0.0010 ~ 0.010% The 7th Honoki is raass%,
C ： 0. 025 ~0. 15% C: 0.025 to 0.15%
Si： 1. 0 %以下、 Si: 1.0% or less,
Mn： 2. 0 %以下、 Mn: 2.0% or less,
P ： 0. 08%以下、 P: 0.08% or less,
S ： 0. 02%以下、 S: 0.02% or less,
A1： 0. 02%以下、 A1: 0.02% or less,
N ： 0. 0050〜0. 0250% N: 0.0050 to 0.0250%
を含み、 かつ N/A1が 0. 3 以上、 固溶状態としての Nを 0. 0010%以上含有し、 残部 が Feおよび不可避的不純物からなる組成と、 平均結晶粒径： 10 m 以下のフェライ ト相を面積率で 80%以上含み、 さらに第 2相として面積率で 2 %以上のマルテンサ ィ ト相を含む組織とを有し、 r値： 1. 2 以上であることを特徴とする高 r値と優れ た歪時効硬化特性および常温非時劾性を有する高張力冷延鋼板である。 And at least 0.3% of N / A1 and 0.0010% of N as a solid solution, with the balance being Fe and unavoidable impurities, and a ferrite with an average grain size of 10 m or less. And a structure containing a martensite phase having an area ratio of 2% or more as a second phase, and having an r value of 1.2 or more. It is a high-tensile cold-rolled steel sheet that has an r-value, excellent strain age hardening characteristics, and non-impeachability at room temperature.
第 7のキ発明では、 前記組成に加えてさらに、 mass%で、 下記 d群〜 g群のうち の 1群または 2群以上を含むことが好ましい。 記  In the seventh invention, it is preferable that, in addition to the above composition, one or more of the following groups d to g are further included in mass%. Record
d群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0 %以下 e群： Nb、 Ti、 Vのうちの 1種または 2種以上を合計で 0. 1 %以下  Group d: One or more of Cu, Ni, Cr, and Mo is 1.0% or less in total. Group e: One or more of Nb, Ti, and V is 0.1 in total. %Less than
f 群： Bを 0. 0030%以下  f group: B is less than 0.0030%
g群： Ca、 REM の 1種または 2種を合計で 0. 0010〜0. 010 %  g group: Ca or REM 1 or 2 kinds in total 0.0010 to 0.010%
第 8の本発明は、 mass %で、 The eighth present invention provides, in mass%,
C ： 0. 025 〜0. 15% C: 0.025 to 0.15%
Si： 1. 0 %以下、 Si: 1.0% or less,
Mn： 2. 0 %以下、 Mn: 2.0% or less,
P ： 0. 08%以下、 P: 0.08% or less,
S ： 0. 02%以下、 S: 0.02% or less,
A1： 0. 02%以下、 A1: 0.02% or less,
N ： 0. 0050〜0. 0250% N: 0.0050 to 0.0250%
を含み、 かつ NZA1が 0. 3 以上である組成の鋼スラブを、 And a steel slab having a composition of which NZA1 is 0.3 or more,
スラブ加熱温度： 1000 以上に加熱し、 Slab heating temperature: Heat to more than 1000,
粗圧延してシートパーとなし、 Rough rolling and sheet par, none
該シートパーに仕上圧延出側温度： 800 °C以上とする仕上圧延を施し、 卷取温度： 650 で以下で卷き取り熱延板とする熱間圧延工程と、 A hot rolling step of subjecting the sheet par to finish rolling at a finish-rolling exit temperature of 800 ° C. or higher, and forming a rolled hot rolled sheet at a winding temperature of 650 or less;
該熱延板に酸洗およぴ冷間圧.延を施し冷延板とする冷間圧延工程と、 Pickling and cold-pressing the hot-rolled sheet; cold-rolling a cold-rolled sheet;
該冷延板に焼鈍温度：再結晶温度以上 800 °C以下で箱焼鈍を施し、ついで焼鈍温度： Ac l変態点〜 （Ac 3変態点一 20°C) で連続焼鈍を行い、 その後 500 ¾以下の温度 域まで冷却速度： 10〜 300で/ sで冷却する冷延板焼鈍工程とを、 順次施すことを特 徴とする r値： 1. 2 以上の高 r値と優れた歪時効硬化特性おょぴ常温非時効性を有 する高張力冷延鋼板の製造方法である。 第 8の本発明では、 前記連続焼鈍後の冷却に引き続いて、 前記冷却の冷却停止温度 以下 350 °C以上の温度域で滞留時間 20 s以上の過時効処理を行うことが好ましい。 第 8の本発明では、 前記組成に加えてさらに、 mass%で、 下記 d群〜 g群のうちの 1群または 2群以上を含むことが好ましい。 Annealing temperature is applied to the cold-rolled sheet at a temperature not lower than the recrystallization temperature and not higher than 800 ° C. A continuous annealing process is performed at the transformation point of Ac1 to (transformation point of Ac3-20 ° C), and thereafter, the cold-rolled sheet annealing process of cooling at a rate of 10 to 300 / s to a temperature range of 500 ° C or less is performed sequentially. R-value characterized by application: This is a method of manufacturing a high-strength cold-rolled steel sheet having a high r-value of 1.2 or more and excellent strain aging hardening properties and non-aging at room temperature. In the eighth aspect of the present invention, it is preferable that, after the cooling after the continuous annealing, an overaging treatment for a residence time of 20 s or more is performed in a temperature range of 350 ° C. or less to the cooling stop temperature of the cooling. In the eighth aspect of the present invention, it is preferable that one or more of the following groups d to g be further contained in mass% in addition to the composition.
記 .  Record .
d群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0 %以下、 Group d: One or more of Cu, Ni, Cr and Mo are 1.0% or less in total,
e群： Nb、 Ti、 Vのうちの 1種または 2種以上を合計で 0. 1 %以下 Group e: 0.1% or less in total of one or more of Nb, Ti, and V
f 群： Bを 0. 0030%以下  f group: B is less than 0.0030%
g群： Ca、 REM の 1種または 2種を合計で 0. 0010〜0. 010 % g group: Ca or REM 1 or 2 kinds in total 0.0010 to 0.010%
図面の簡単な説明 BRIEF DESCRIPTION OF THE FIGURES
図 1 鋼成分 （N% - 14/93 - Nb% - 14/27 - Al% - 14/11 - B %) と成形後引張 強度上昇代 （ATS) との関係を示したグラフである。 Figure 1 is a graph showing the relationship between the steel composition (N%-14/93-Nb%-14/27-Al%-14/11-B%) and the post-forming tensile strength rise (ATS).
図 2 Nb， B複合添加鋼における B含有量と ATSとの関係を示したグラフである。 図 3 固溶 Cが多い鋼 B (従来鋼） と固溶 Nが多い鋼 A (発明鋼） において、 低温 温度域での成形後熱処理による引張強度上昇代の違いを比較して示したグラフであ る。 Fig. 2 is a graph showing the relationship between B content and ATS in Nb and B composite added steel. Fig. 3 is a graph comparing the difference in tensile strength rise between steel B with a large amount of solid solution C (conventional steel) and steel A with a large amount of solid solution N (inventive steel) due to post-forming heat treatment at a low temperature range. is there.
図 4 常温時効による伸びの低下量 （ΔΕ1) と成形後引張強度上昇代 （ATS) に及 ぼす、 結晶粒径 dと鋼成分 （N% - 14/93 · Nb% - 14/27 · Al% - 14/11 · B %) との影響を示したグラフである。 Figure 4 Grain size d and steel composition (N%-14/93 · Nb%-14/27 · Al) affecting the amount of decrease in elongation due to normal temperature aging (ΔΕ1) and increase in tensile strength after forming (ATS) %-14/11 · B%).
図 5 TSX i fi、 BH、 ATSと N/(A1+Nb+B) との関係を示すグラフである。 図 6 TSX r値、 ΒΗ、 Δ T Sと B量との関係を示すグラフである。 Figure 5 is a graph showing the relationship between TSX i fi, BH, ATS and N / (A1 + Nb + B). FIG. 6 is a graph showing the relationship between the TSX r value, ΒΗ, ΔTS, and the B amount.
発明を実施するための最良の形態 BEST MODE FOR CARRYING OUT THE INVENTION
第 1の本発明において鋼板の成分組成を前記の範囲に限定した理由について説明 する。  The reason for limiting the component composition of the steel sheet to the above range in the first invention will be described.
C ： 0. 01masso/o未満 C: less than 0.01 mass o / o
Cは、 できるだけ少量であるほど深絞り性に優れ、 プレス成形性の面で有利であ る。 また、 冷間圧延後の焼鈍過程において Nb Cの再溶解が進行し結晶粒内の固溶 C が増加して、 耐常温時効性の低下を招き易い。 従って、 C量は 0. 01mass%未満に抑 制することが好ましい。より好ましくは 0， 0050mass%以下であり、より好ましくは 0. 0030mass%以下である。  The smaller the amount of C, the better the deep drawability and the more advantageous in terms of press formability. Also, during the annealing process after cold rolling, the re-dissolution of NbC proceeds, solute C in the crystal grains increases, and the normal-temperature aging resistance tends to decrease. Therefore, the amount of C is preferably suppressed to less than 0.01 mass%. It is more preferably at most 0.0050 mass%, more preferably at most 0.0030 mass%.
Si： 0. 005 ~ 1. 0 mass% Si: 0.005 to 1.0 mass%
Siは、 伸びの低下を抑制し、 また強度を向上させる有用成分であるが、 含有量が 0. 005mass%に満たないとその添加効果に乏しく、 一方 1. 0mass%を超えると表面性 状を悪化させ、 延性の低下を招くので、 Siは 0. 005~1. 0 mass%の範囲に限定した。 より好ましくは 0. 01〜0. 75mass%の範囲である。  Si is a useful component that suppresses the decrease in elongation and improves the strength.However, if the content is less than 0.005 mass%, the effect of its addition is poor, and if it exceeds 1.0 mass%, the surface properties will be deteriorated. Therefore, Si was limited to the range of 0.005 to 1.0 mass% because it deteriorated the ductility. More preferably, it is in the range of 0.01 to 0.75 mass%.
Mn： 0. 01~ 1. 5 mass% Mn: 0.01 to 1.5 mass%
Mnは、 鋼の強化成分として有用なだけでなく、 Mn Sを形成して Sによる脆化を抑 制する作用があるが、 含有量が 0. 01mass%に満たないとその添加効果に乏しく、 一 方 1. 5 mass %を超えると表面性状の悪化や延性の低下を招くので、 Mnは 0. 01〜1. 5 m ass%の範囲で含有させるものとした。 より好ましくは 0. 10〜0. 75mass%である。 P ： 0. 10mass%以下  Mn is not only useful as a strengthening component of steel, but also has the effect of forming Mn S to suppress embrittlement due to S. However, if the content is less than 0.01 mass%, the effect of its addition is poor. On the other hand, if the content exceeds 1.5 mass%, the surface properties deteriorate and the ductility decreases. Therefore, Mn is contained in the range of 0.01 to 1.5 mass%. More preferably, it is 0.10 to 0.75 mass%. P: 0.10 mass% or less
Pは、 固溶強化成分として鋼の強化に有効に寄与するが、 0. 10mass%を超えて添 加すると、 （FeNb) xPなどのリ ン化物を形成するため深絞り性が低下する。 従って、 Pは 0. 10mass%以下に限定した。  P effectively contributes to the strengthening of steel as a solid solution strengthening component, but when added in excess of 0.10 mass%, deep drawability is reduced due to the formation of phosphides such as (FeNb) xP. Therefore, P was limited to 0.10 mass% or less.
S ： 0. 01mass%以下 S: 0.01 mass% or less
Sが多量に含有されると介在物量が増大し、 延性の低下を招くので、 Sの混入は 極力避けることが望ましいが、 0. 01mass%までは許容される。 Al： 0. 005 ~0. 030 mass% If a large amount of S is contained, the amount of inclusions increases and the ductility is reduced. Therefore, it is desirable to avoid the incorporation of S as much as possible, but up to 0.01 mass% is allowable. Al: 0.005 ~ 0.030 mass%
Alは、 脱酸剤として、 また炭窒化物形成成分の歩留りを向上するために添加する が、 含有量が 0. 005mass%未満では十分な効果がなく、 一方 0. 030raass%を超える 添加は鋼中に添加すべき N量の増大を招き、 製鋼時のスラブ欠陥が発生し易くなる。 従って、 A1は 0. 005〜0. 030 mass%の範囲で含有させるものとした。  Al is added as a deoxidizing agent and to improve the yield of carbonitride-forming components.However, if the content is less than 0.005 mass%, it has no sufficient effect, whereas if it exceeds 0.003 mass%, steel is added. This leads to an increase in the amount of N to be added into the steel, which tends to cause slab defects during steelmaking. Therefore, A1 is contained in the range of 0.005 to 0.30 mass%.
N： 0. 005 〜0. 040 mass% N: 0.005 to 0.440 mass%
Nは、 本発明において、 歪み時効硬化特性を鋼板に付与する役割を果たす重要な 元素である。 しかしながら、 含有量が 0. 005mass%に満たないと十分な歪み時効硬 化特性が得られず、 一方 0. 040mass%を超える多量添加はプレス成形性の低下を招 く。 従って、 Nは 0. 005〜0. 040 mass%の範囲で含有させるものとした。 なお好ま しくは 0. 008~0. 015mass%である。  N is an important element that plays a role in imparting strain age hardening characteristics to a steel sheet in the present invention. However, if the content is less than 0.005 mass%, sufficient strain aging hardening properties cannot be obtained, while addition of a large amount exceeding 0.004 mass% causes a decrease in press formability. Therefore, N was contained in the range of 0.005 to 0.440 mass%. It is more preferably 0.008 to 0.015 mass%.
B ： 0. 0001〜0. 003mass% B: 0.0001 to 0.003 mass%
Bは、 Nbと複合添加することにより、 熱延組織およぴ冷延再結晶組織を効果的に 微細化し、 また耐二次加工脆性を改善する作用がある。 しかしながら、 含有量 が 0. 0001mass%未満では十分な微細化効果が得られず、一方 0. 003mass%を超えると B N析出量が増大するだけでなく、 スラブ加熱段階での溶体化に支障を来すように なる。 従って、 Bは 0. 0001〜0. 003mass%の範囲で含有させるものとした。 なお好ま しくは 0. 0001~0. 0015mass%で、 より好ましくは 0. 0007〜0. 0012raass%である。 Nb： 0. 005 〜0· 050 mass%  B, when added in combination with Nb, has the effect of effectively refining the hot-rolled structure and the cold-rolled recrystallized structure and improving the secondary work brittleness resistance. However, if the content is less than 0.0001 mass%, a sufficient refining effect cannot be obtained.On the other hand, if the content exceeds 0.003 mass%, not only does the amount of BN precipitate increase, but also hinders solution formation in the slab heating step. It will be. Therefore, B is contained in the range of 0.0001 to 0.003 mass%. It is more preferably 0.0001 to 0.0015 mass%, and more preferably 0.0007 to 0.0012raass%. Nb: 0.005 to 050 mass%
Nbは、 Bとの複合添加によって熱延組織およぴ冷延再結晶焼鈍組織の微細化に有 効に寄与し、 また固溶 Cを Nb Cとして固定する作用がある。 さらに、 Nbは NbNとい つた窒化物を形成し、 冷延再結晶焼鈍組織の微細化に寄与する。 しかしながら、 Nb 量が 0. 005raass %に満たないと固溶 Cを析出固定することが困難となるばかりでな く、熱延組織およぴ冷延再結晶焼鈍組織の微細化が不十分となり、一方 0. 050 mass% を超えると延性の低下を招く。 従って、 Nbは 0. 005〜0. 050 mass%範囲で含有させ るものとした。 好ましくは 0. 010〜0. 030 raass%である。 また、 上述したとおり、 Nbは、 固溶 Cを Nb Cとして固定する作用がある。 また、 N bNといった窒化物を形成する。 同様に、 A1および Bはそれぞれ A1N, B Nを形成す る。 従って、 崮溶 N量を十分に確保すると共に、 固溶 Cを十分に低減するためには、 次式（1) , (2)の関係を満足させることが重要である。 Nb contributes effectively to the refinement of the hot-rolled structure and the cold-rolled recrystallization-annealed structure by the combined addition with B, and has the effect of fixing solid solution C as NbC. In addition, Nb forms a nitride called NbN, which contributes to the refinement of the cold-rolled recrystallization annealing structure. However, if the Nb content is less than 0.005 raass%, it becomes difficult not only to precipitate and fix solid solution C, but also to make the hot-rolled structure and the cold-rolled recrystallization annealed structure insufficiently refined. On the other hand, if it exceeds 0.050 mass%, ductility is reduced. Therefore, Nb is contained in the range of 0.005 to 0.050 mass%. Preferably, it is 0.010 to 0.030 raass%. Further, as described above, Nb has an effect of fixing solid solution C as NbC. Also, a nitride such as NbN is formed. Similarly, A1 and B form A1N and BN, respectively. Therefore, it is important to satisfy the relations of the following equations (1) and (2) in order to sufficiently secure the dissolved N amount and sufficiently reduce the solute C.
N %≥0. 0015 + 14/93 · Nb% + 14/27 · Al% + 14/11 · B % --- (1) C %≤ 0. 5 · (12/93) - Nb% -― (2) また、 この発明において、 高い歪時効特性を得ると共に、 時効劣化を防止するた めには、 結晶粒径を小さくすることが好適である。  N% ≥0.0015 + 14 / 93Nb% + 14 / 27Al% + 14 / 11B% --- (1) C% ≤0.55 (12/93)-Nb%- (2) In the present invention, in order to obtain high strain aging characteristics and prevent aging deterioration, it is preferable to reduce the crystal grain size.
すなわち、 前掲図 4に示したように、 結晶粒径 dを 20 /z ra 以下まで小さくするこ とによって、 （N % — 14/93 · Nb% - 14/27 · Al% - 14/11 · B %) ≥0. 0015mass% と比較的多量の固溶 Nを含有する場合においても、 Δ Ε1を 2. 0%以下まで抑制する ことが可能となる。 なお、 より好適には、 結晶粒径 dを 15 μ ιη 以下まで小さくする ことが好ましい。 というのは、 図 4に示したように、 結晶粒径 dを 以下まで 小さくすると、 Δ Ε1を 1. 5%以下まで抑制することが可能となるからである。 第 2の本発明による製造条件について述べる。  In other words, as shown in Figure 4 above, by reducing the crystal grain size d to 20 / z ra or less, (N% — 14/93 · Nb%-14/27 · Al%-14/11 · B%) ≥0.0015 mass%, it is possible to suppress ΔΕ1 to 2.0% or less even when a relatively large amount of solute N is contained. It is more preferable to reduce the crystal grain size d to 15 μιη or less. This is because, as shown in FIG. 4, when the crystal grain size d is reduced to below, Δ 以下 1 can be suppressed to 1.5% or below. The manufacturing conditions according to the second invention will be described.
上記の好適成分組成になる鋼を、 転炉等の公知の溶製方法で溶製し、 造塊法または 連続铸造法で鋼片とする。 Steel having the above-mentioned preferred composition is smelted by a known smelting method such as a converter, and is made into a slab by an ingot-making method or a continuous sintering method.
ついで、 この鋼片を、 加熱、 均熱したのち、 熱間圧延を施して熱延板とする。 こ の発明では、 熱間圧延の加熱温度は特に規定するものではないが、 深絞り性の向上 のためには固溶 Cを固定し炭化物として析出させておくのが有利であり、 このため には熱間圧延の加熱温度は 130(TC以下にするのが好ましい。 また、 加工性のより一 層の向上のためには 1150°C以下とするのが好ましい。 しかしながら、 加熱温度が 90 0°C未満では、 加工性の改善は飽和し、 逆に熱間圧延時の圧延負荷が増大して圧延ト ラブルが発生する危険性が増大するので、 加熱温度の下限は 900 とするのが好ま しい。 次に、 熱間圧延における全圧下率は 70%以上とすることが好ましい。 というのは、 全圧下率が 70%未満では熱延板の結晶粒微細化が不十分となるからである。 Next, the steel slab is heated and soaked, and then hot-rolled to obtain a hot-rolled sheet. In the present invention, the heating temperature of hot rolling is not particularly specified, but it is advantageous to fix solid solution C and precipitate as carbide in order to improve the deep drawability. The hot rolling temperature is preferably set to 130 (TC or lower. In order to further improve the workability, the temperature is preferably set to 1150 ° C. or lower. However, the heating temperature is 900 ° C. If the temperature is less than C, the improvement in workability is saturated, and conversely, the rolling load during hot rolling increases and the risk of rolling trouble increases, so the lower limit of the heating temperature is preferably 900. . Next, the total reduction in hot rolling is preferably set to 70% or more. This is because if the total draft is less than 70%, the grain refinement of the hot-rolled sheet will be insufficient.
また、 熱間圧延における仕上圧延は 960~650 ¾の温度域で終了するのが好ましく、 熱間圧延仕上温度は、 Ar 3変態点以上の y域であっても、 Ars変態点以下の α域で あってもよい。 熱間圧延仕上温度が ％0 超えると熱延板の結晶粒が粗大化し、 冷 延 ·焼鈍後の深絞り性が劣化する。 一方 650 未満では、 変形抵抗が増加するため 熱延負荷の増大を招き圧延が困難になる。 Further, the finish rolling in the hot rolling is preferably completed in a temperature range of 960 to 650 ° C, and the hot rolling finish temperature may be in the y range above the Ar 3 transformation point but in the α range below the Ars transformation point. It may be. If the hot rolling finish temperature exceeds 0%, the crystal grains of the hot rolled sheet become coarse, and the deep drawability after cold rolling and annealing deteriorates. On the other hand, if it is less than 650, the deformation resistance increases, so that the hot rolling load increases and rolling becomes difficult.
上記の熱間仕上圧延終了後は、 直ちに冷却を開始することによって、 正常粒成長 を防止すると共に、 冷却過程での A1N析出を抑制することが望ましい。  It is desirable to start cooling immediately after the above hot finish rolling to prevent normal grain growth and to suppress A1N precipitation during the cooling process.
ここに、 上記の冷却処理条件については特に限定するものではないが、 冷却開始 時間は、 仕上圧延終了後、 好ましくは 1. 5秒以内、 より好ましくは 1. 0秒以内、 さ らに好ましくは 0. 5秒以内とすることが望ましい。 というのは、 圧延終了後直ちに 冷却すると、 歪が蓄積した状態での過冷度が大きくなるため、 より多くのフェライ ト核が生成し、 フェライ ト変態が促進されると共に、 γ相中の固溶 Νがフェライ ト 粒内に拡散するのが抑制され、 フェライ ト粒界に存在する固溶 Ν量が増加するから である。  Here, the above cooling treatment conditions are not particularly limited, but the cooling start time is preferably within 1.5 seconds, more preferably within 1.0 seconds, and more preferably after the finish rolling. It is desirable to keep it within 0.5 seconds. This is because cooling immediately after the end of rolling increases the degree of supercooling with accumulated strain, so that more ferrite nuclei are generated, ferrite transformation is accelerated, and solidification in the γ phase is promoted. This is because diffusion of the melt into the ferrite grains is suppressed, and the amount of solid solution present at the ferrite grain boundaries increases.
また、 冷却速度については、 固溶 Νを確保するために、 10°C/s以上とするのが好 ましい。 なお、 特に熱延仕上温度が Ar 3変態点以上の場合には、 冷却速度を 50°C/s 以上とすることが、 固溶 Nを確保する上でより好適である。 The cooling rate is preferably set to 10 ° C / s or more in order to secure solid solution. In particular, when the hot-rolling finishing temperature is equal to or higher than the Ar 3 transformation point, it is more preferable to set the cooling rate to 50 ° C / s or higher in order to secure solid solution N.
ついで、 熱延板をコイルに卷き取る。 この卷取り温度は高温ほど炭化物の粗大化 には有利であるが、 800°Cを超えると熱延板表面に形成されるスケールが厚くなつ てスケール除去作業の負荷が增大するだけでなく、 窒化物形成が進行しコイル長手 方向の固溶 N量の変動を招き、 一方卷取り温度が 400eC未満では、 卷取り作業が困 難になるので、 熱延板の卷取り温度は 800〜400 °Cの範囲とする必要がある。 Next, the hot rolled sheet is wound up into a coil. The higher the winding temperature, the more advantageous is the coarsening of carbides. However, if it exceeds 800 ° C, the scale formed on the surface of the hot-rolled sheet becomes thicker, which not only increases the load of scale removal work, but also increases the load. nitridation proceeds leads to variation in the coil longitudinal direction of solute N amount, while in卷取Ri temperature is less than 400 e C, since the flame Certificates up work piece,卷取Ri temperature of the hot-rolled sheet is 800 It must be in the range of 400 ° C.
ついで、 熱延板に冷間圧延を施すが、 かかる冷間圧延における圧下率は 60〜95% とする必要がある。 というのは、 冷間圧延の圧下率が 60%未満では高い r値が期待 できず、 一方 95%を超えると r値がかえって低下するからである。 Next, the hot rolled sheet is subjected to cold rolling, and the rolling reduction in such cold rolling needs to be 60 to 95%. This is because a high r value is expected when the rolling reduction of cold rolling is less than 60%. No, on the other hand, if it exceeds 95%, the r-value will decrease.
上記のような冷間圧延を施された冷延板は、 次に再結晶焼鈍に供される。 焼鈍方 法は、 連続焼鈍であっても、 パッチ焼鈍であっても何れでも良いが、 連続焼鈍の方 が有利である。 なお、 この連続焼鈍は、 通常の連続焼鈍ラインにおける処理あるい は連続溶融亜鉛めつきラインにおける処理の何れであっても良い。  The cold rolled sheet that has been subjected to the above cold rolling is then subjected to recrystallization annealing. The annealing method may be either continuous annealing or patch annealing, but continuous annealing is more advantageous. The continuous annealing may be either a normal continuous annealing line or a continuous hot-dip galvanizing line.
また、 焼鈍条件は 650t以上、 5秒以上とすることが好ましい。 というのは、 焼 鈍温度が 650 C未満、 焼鈍条件が 5秒未満では再結晶が完了せず、 そのため深絞り 性が低下するからである。 深絞り性をより向上させるためには、 800で以上のフエ ライ ト単相域で 5秒以上焼鈍することが望ましい。  The annealing conditions are preferably 650 t or more and 5 seconds or more. This is because if the annealing temperature is less than 650 C and the annealing condition is less than 5 seconds, recrystallization is not completed, and the deep drawability decreases. In order to further improve the deep drawability, it is desirable to perform annealing in a ferrite single phase region of 800 or more for 5 seconds or more.
またより高温の α + y二相域での焼鈍により部分的に o:→ y変態が生じることで { 1 1 1 } 集合組織が発達し r値が向上するが、 a→y変態が完全に進行した場合 は集合組織がランダム化するため、 r値が低下し深絞り性が損なわれる。 Also, annealing in the α + y two-phase region at a higher temperature partially causes the o: → y transformation to develop {1 1 1} texture and improve the r value, but the a → y transformation is completely If it progresses, the texture will be randomized, so the r-value will decrease and the deep drawability will be impaired.
なお、 焼鈍温度の上限は 900 とすることが好ましい。 というのは、 焼鈍温度が 90 0でを超えると、 炭化物の再溶解が進行し固溶 Cが過度に増加するため、 遅時効性が 低下するからであり、 また α— γ変態が生じた場合は集合組織がランダム化するた め、 r値が低下し深絞り性が損なわれるからである。 Preferably, the upper limit of the annealing temperature is 900. This is because, when the annealing temperature exceeds 900, the re-dissolution of the carbide proceeds and the solid solution C excessively increases, so that the delayed aging property decreases, and when the α-γ transformation occurs This is because the texture is randomized, so that the r-value decreases and the deep drawability is impaired.
さらに、 上記した再結晶焼鈍における昇温過程において、 500°Cから再結晶温度 までの温度域を徐熱とし、 A1 N等を十分に析出させることによって、 鋼板の結晶粒 径を効果的に小さくすることができる。  Furthermore, during the temperature rise process in the above-mentioned recrystallization annealing, the temperature range from 500 ° C to the recrystallization temperature is gradually heated, and A1N etc. are sufficiently precipitated to effectively reduce the crystal grain size of the steel sheet. can do.
ここに、 上記したような制御加熱を施すべき温度域は、 A1N等が析出し始める 500 °Cから再結晶温度までとする。  Here, the temperature range in which the above-described controlled heating should be performed is from 500 ° C. at which A1N or the like starts to precipitate to the recrystallization temperature.
また、 昇温速度は 1〜20°C/sの範囲とすることが好ましい。 というのは、 昇温速 度が 20¾ 超では十分な析出量が得られず、 一方 l °C/s未満では析出物が粗大化し て粒成長の抑制効果が弱まるからである。  Further, the heating rate is preferably in the range of 1 to 20 ° C / s. This is because if the heating rate is higher than 20 ° C, a sufficient amount of precipitate cannot be obtained, while if it is lower than l ° C / s, the precipitate becomes coarse and the effect of suppressing the grain growth is weakened.
なお、 上記のような再結晶焼鈍後に、 さらに形状矯正、表面粗さ調整のため、 10% 以下の調質圧延を行ってもよい。 また再結晶焼鈍における均熱後の冷却速度は 1 0 ~ 5 °C / s とするのが好まし い。 というのは冷却速度が 1 0 °CZ s以下では冷却中に粒成長を起こし結晶粒の粗 大化がおこり歪時効特性、 および常温での時効特性が低下する。 一方、 5 0 °CZ s 以上では固溶状態の Nの粒界への拡散が十分には起こらず、 常温での時効特性を低 下させる。 なお好ましくは 1 0〜3 O tZ sである。 After the recrystallization annealing as described above, a temper rolling of 10% or less may be performed for further shape correction and surface roughness adjustment. Further, the cooling rate after soaking in the recrystallization annealing is preferably set to 10 to 5 ° C / s. At a cooling rate of 10 ° CZs or less, grain growth occurs during cooling, crystal grains become coarse, and strain aging characteristics and aging characteristics at room temperature decrease. On the other hand, above 50 ° CZ s, diffusion of N in the solid solution state to the grain boundary does not occur sufficiently, and the aging characteristics at room temperature deteriorate. Still more preferably, it is 10 to 3 OtZs.
上記の再結晶焼鈍に引き続き、 必要に応じて、 溶融亜鉛めつき処理ついで加熱合 金化処理を行うことにより合金化溶融亜鉛めつき鋼板とする。  Subsequent to the recrystallization annealing described above, if necessary, a hot-dip galvanizing treatment followed by a heat alloying treatment is performed to form an alloyed hot-dip galvanized steel sheet.
かかる溶融亜鉛めつき処理おょぴ合金化処理については特に限定されることはな く、 従来公知の方法に従って行えば良い。  There is no particular limitation on the hot-dip galvanizing treatment and alloying treatment, and the treatment may be performed according to a conventionally known method.
なお、 合金化溶融亜鉛めつき鋼板としたのち、 加工性の向上や加工後の外観向上 のために調質圧延を施した鋼板 （ダル仕上鋼板、 ブライ ト仕上鋼板、 表面に特定の 粗度パターンを形成した鋼板） 、 表面に防鲭油、 潤滑油などの油胰層を有する鋼板 など、 通常に薄鋼板として採用する表面処理を施した鋼板において、 この発明の成 分範囲であればこの発明の効果を十分に享受できる。  In addition, steel sheets with alloyed hot-dip galvanized steel and then temper-rolled to improve workability and appearance after processing (dull-finished steel sheet, bright-finished steel sheet, and a specific roughness pattern on the surface) Steel sheet which has been subjected to surface treatment that is usually used as a thin steel sheet, such as a steel sheet having an oil layer such as a fireproof oil or a lubricating oil on the surface. Effect can be fully enjoyed.
かくして、 優れた深絞り性を有するだけでなく、 プレス成形一熱処理により引張 強度が増加する、 歪み時効硬化特性に優れた冷延鋼板おょぴ合金化溶融亜鉛めつき 鋼板を得ることができる。 第 3の本発明において鋼板の成分組成を前記の範囲に限定した理由について説明 する。  Thus, it is possible to obtain a cold-rolled steel sheet and an alloyed hot-dip galvanized steel sheet having not only an excellent deep drawability but also an increase in tensile strength by press forming and heat treatment, and an excellent strain aging hardening property. The reason for limiting the component composition of the steel sheet to the above range in the third invention will be described.
C ： 0. 01mass%未満  C: less than 0.01 mass%
Cは、 できるだけ少量であるほど深絞り性に優れ、 プレス成形性の面で有利であ る。 また、 冷間圧延後の焼鈍過程において Nb Cの再溶解が進行し結晶粒内の固溶 C が増加して、 耐常温時効性の低下を招き易い。 従って、 C量は 0. 01mass%未満に抑 制することが好ましい。より好ましくは 0. 0050raass%以下であり、より好ましくは 0. 0030raass%以下である。 なお、 強度確保と結晶粒粗大化防止の観点からは、 Cは 0. 0 005%以上含有させるのが望ましい。 The smaller the amount of C, the better the deep drawability and the more advantageous in terms of press formability. Also, during the annealing process after cold rolling, the re-dissolution of NbC proceeds, solute C in the crystal grains increases, and the normal-temperature aging resistance tends to decrease. Therefore, the amount of C is preferably suppressed to less than 0.01 mass%. More preferably, it is not more than 0.0050raass%, more preferably not more than 0.0030raass%. From the viewpoint of securing strength and preventing coarsening of crystal grains, C is 0.0. It is desirable to contain 005% or more.
Si： 0. 005 〜1. 0 mass% Si: 0.005 to 1.0 mass%
Siは、 伸びの低下を抑制し、 また強度を向上させる有用成分であるが、 含有量が 0. 005mass%に満たないとその添加効果に乏しく、一方 1. 0mass%を超えると表面性 状を悪化させ、 延性の低下を招くので、 Siは 0. 005~1. 0 mass%の範囲に限定した。 より好ましくは 0. 01〜 75mass%の範囲である。  Si is a useful component that suppresses the decrease in elongation and improves the strength.However, if the content is less than 0.005 mass%, the effect of its addition is poor, while if it exceeds 1.0 mass%, the surface properties will deteriorate. Therefore, Si was limited to the range of 0.005 to 1.0 mass% because it deteriorated the ductility. More preferably, it is in the range of 0.01 to 75 mass%.
Mn： 0. 01〜1. 5 raass % Mn: 0.01 to 1.5 raass%
Mnは、 鋼の強化成分として有用なだけでなく、 Mn Sを形成して Sによる脆化を抑 制する作用があるが、 含有量が 0. 01mass%に満たないとその添加効果に乏しく、 一 方 1. 5 mass%を超えると表面性状の悪化や延性の低下を招くので、 Mnは 0. 01~1. 5 m ass%の範囲で含有させるものとした。 より好ましくは 0. 10~0. 75raass%である。  Mn is not only useful as a strengthening component of steel, but also has the effect of forming Mn S to suppress embrittlement due to S. However, if the content is less than 0.01 mass%, the effect of its addition is poor. On the other hand, if the content exceeds 1.5 mass%, deterioration of the surface properties and decrease in ductility are caused. Therefore, Mn is contained in the range of 0.01 to 1.5 mass%. More preferably, it is 0.10 to 0.75raass%.
P ： 0. 10nmss%以下 P: 0.10 nmss% or less
Pは、 固溶強化成分として鋼の強化に有効に寄与するが、 0. 10raass%を超えて添 加すると、 （FeNb) xPなどのリン化物を形成するため深絞り性が低下する。 従って、 Pは 0. 10mass%以下に限定した。  P effectively contributes to the strengthening of steel as a solid solution strengthening component. However, if added in excess of 0.10 raass%, phosphides such as (FeNb) xP are formed, and the deep drawability decreases. Therefore, P was limited to 0.10 mass% or less.
S ： 0. 01niass%以下 S: 0.01 niass% or less
Sが多量に含有されると介在物量が増大し、 延性の低下を招くので、 Sの混入は 極力避けることが望ましいが、 0. 01mass%までは許容される。  If a large amount of S is contained, the amount of inclusions increases and the ductility is reduced. Therefore, it is desirable to avoid the incorporation of S as much as possible, but up to 0.01 mass% is allowable.
A1 ： 0. 005 〜0. 030 mass% A1: 0.005 to 0.30 mass%
A1は、 脱酸剤として、 また炭窒化物形成成分の歩留りを向上するために添加する が、 含有量が 0. 005mass%未満では十分な効果がなく、 一方 0. 030mass%を超える 添加は鋼中に添加すべき N量の増大を招き、製鋼時のスラブ欠陥が発生し易くなる。 従って、 A1は 0. 005~0. 030 mass%の範囲で含有させるものとした。  A1 is added as a deoxidizing agent and to improve the yield of carbonitride-forming components.However, if the content is less than 0.005% by mass, it has no sufficient effect, whereas if it exceeds 0.003% by mass, the addition of A1 This leads to an increase in the amount of N to be added into the steel, which tends to cause slab defects during steelmaking. Therefore, A1 is contained in the range of 0.005 to 0.30 mass%.
N ： 0. 005 〜 040 mass % N: 0.005 to 040 mass%
Nは、 本発明において、 歪み時効硬化特性を鋼板に付与する役割を果たす重要な 元素である。 しかしながら、 含有量が 0. 005mass%に満たないと十分な歪み時効硬 化特性が得られず、 一方 0.040mass%を超える多量添加はプレス成形性の低下を招 く。 従って、 Nは 0.005〜0.040 mass%の範囲で含有させるものとした。 なお好ま しくは 0.008〜0.015mass%である。 N is an important element that plays a role in imparting strain age hardening characteristics to a steel sheet in the present invention. However, if the content is less than 0.005 mass%, sufficient strain aging However, the addition of a large amount exceeding 0.040 mass% leads to a decrease in press formability. Therefore, N was contained in the range of 0.005 to 0.040 mass%. Preferably, it is 0.008 to 0.015 mass%.
B： 0.0001〜0.003mass% B: 0.0001-0.003mass%
Bは、 Nbと複合添加することにより、 熱延組織およぴ冷延再結晶組織を効果的に 微細化し、 また耐二次加工脆性を改善する作用がある。 しかしながら、 含有量 が 0.0001mass%未満では十分な微細化効果が得られず、 一方 0.003mass%を超えると BN析出量が増大するだけでなく、 スラブ加熱段階での溶体化に支障を来すように なる。 従って、 Bは 0.0001〜0.003raass%の範囲で含有させるものとした。 なお好ま しくは 0.0001〜0.0015mass%で、 より好ましくは 0.0007〜0· 0012mass%である。 Nb： 0.005 〜0·050 %、 Ti： 0.005 〜0.070 %、 V： 0.005〜0.10%  B, when added in combination with Nb, has the effect of effectively refining the hot-rolled structure and the cold-rolled recrystallized structure and improving the secondary work brittleness resistance. However, if the content is less than 0.0001 mass%, sufficient refining effect cannot be obtained.On the other hand, if the content exceeds 0.003 mass%, not only the BN precipitation amount increases, but also the solution solution at the slab heating stage may be hindered. become. Therefore, B was contained in the range of 0.0001 to 0.003 raass%. It is more preferably 0.0001 to 0.0015 mass%, more preferably 0.0007 to 0.0012 mass%. Nb: 0.005 to 0.050%, Ti: 0.005 to 0.070%, V: 0.005 to 0.10%
Nb, Ti, Vは、 Bと複合添加することにより、 熱延組織およぴ冷延再結晶組織の 微細化に寄与し、 かつ固溶 Cを NbC，TiC,VCとして析出させる働きを有するので必要 に応じて Bとともに添加されるが、 各々 0.005 %未満ではその働きが不十分である。 一方、 Nbは 0.050 %超、 Tiは 0.070 %超、 Vは 0.10%超では延性の劣化を招く。 よ つて Nbは 0.005 〜0.050 %、 Tiは 0.005〜0.070 %、 Vは 0.005 〜0.10%とした。 また、 上述したとおり、 Nbは、 固溶 Cを NbCとして固定する作用がある。 また、 N bNといった窒化物を形成する。 同様に、 A1および Bはそれぞれ A1N， BNを形成す る。 従って、 固溶 N量を十分に確保すると共に、 固溶 Cを十分に低減するためには、 次式（1)， （2)の関係を満足させることが重要である。  Nb, Ti, and V contribute to the refinement of the hot-rolled structure and the cold-rolled recrystallized structure when combined with B, and have the function of precipitating solid solution C as NbC, TiC, and VC. If necessary, they are added together with B, but if each is less than 0.005%, their function is insufficient. On the other hand, if Nb exceeds 0.050%, Ti exceeds 0.070%, and V exceeds 0.10%, ductility deteriorates. Therefore, Nb was 0.005 to 0.050%, Ti was 0.005 to 0.070%, and V was 0.005 to 0.10%. Further, as described above, Nb has an effect of fixing solid solution C as NbC. Also, a nitride such as NbN is formed. Similarly, A1 and B form A1N and BN, respectively. Therefore, it is important to satisfy the following formulas (1) and (2) in order to secure a sufficient amount of solute N and sufficiently reduce the amount of solute C.
N%≥0.0015 + 14/93 - Nb% + 14/27 - Al% + 14/11 · B% —- (1) C%≤ 0.5 · (12/93) · Nb% —-- (2) N% ≥0.0015 + 14/93-Nb% + 14/27-Al% + 14 / 11B% --- (1) C% ≤ 0.5 · (12/93) Nb% ---- (2)
N/A1または N/(A1+Nb+Ti+V+B) ： 0.30以上 N / A1 or N / (A1 + Nb + Ti + V + B): 0.30 or more
A1は A1N を形成して固溶 Nを減らす。固溶 Nの適正量を確保するためには N/A1を 0. 30以上とする必要がある。 また、 Nb，Ti, Vあるいは Bを複合添加する場合は、 これ らも夫々 NbN,TiN, VN，BN を形成して固溶 Nを減らすので、 固溶 Nの適正量を確保す るためには (Al+Nb+Ti+V+B)を 0. 30以上とする必要がある。 A1 forms A1N to reduce solute N. In order to secure an appropriate amount of solid solution N, N / A1 must be 0.30 or more. When Nb, Ti, V or B is added in a complex manner, NbN, TiN, VN and BN are formed to reduce the amount of solute N, so that the appropriate amount of solute N should be secured. For this purpose, (Al + Nb + Ti + V + B) must be 0.30 or more.
固溶 N： 0. 0010%以上 Solid solution N: 0.0010% or more
鋼板の歪時効硬化特性を高めるためには固溶 Nが 0. 0010%以上の含有量で存在す る必要がある。  In order to enhance the strain age hardening characteristics of the steel sheet, it is necessary that solid solution N is present at a content of 0.0010% or more.
ここで、 固溶 N量は、 鋼中の全 N量から析出 N量を差し引いて求めるものとする。 なお、 析出 N量の分析法としては、 本発明者らが種々の分析法を比較検討した結果 によれば、 定電位電解法を用いた電解抽出分析法により求めるのが有効である。 な お抽出分析に用いる地鉄を溶解する方法として、 酸分解法、 ハロゲン法おょぴ電解 法がある。 この中で、 電解法は炭化物、 窒化物などの極めて不安定な析出物を分解 させることなく、 安定して地鉄のみを溶解できる。 電解液としてはァセチル ·ァセ トン系を用いて、 定電位にて電解する。 本発明では定電位電解法を用いて析出 N量 を測定した結果が、 実際の部品強度と最もよい対応を示した。  Here, the amount of solute N is determined by subtracting the amount of precipitated N from the total amount of N in the steel. As a method of analyzing the amount of precipitated N, according to the results of comparative studies of various analysis methods by the present inventors, it is effective to determine the amount by the electrolytic extraction analysis method using the constant potential electrolysis method. There are acid decomposition method, halogen method and electrolysis method to dissolve the iron base used for extraction analysis. Among these, the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides. Electrolysis is performed at a constant potential using an acetyl-aceton system as an electrolyte. In the present invention, the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed the best correspondence with the actual component strength.
このようなことから、 本発明では、 定電位電解法により抽出した残渣を化学分析 して残渣中の N量を求め、 これを析出 N量とする。  For this reason, in the present invention, the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N.
なお、 より高い B Hおよび A T Sを得るためには、 固溶 N量は 0. 0015%以上が好 ましく、 0. 0020%以上がより好ましく、 0. 0030%以上がさらに一層好ましい。  In order to obtain higher BH and ATS, the amount of solid solution N is preferably 0.0015% or more, more preferably 0.0020% or more, and still more preferably 0.0030% or more.
本発明の冷延鋼板は、 上記の組成を有するとともに T S X r値≥
Figure imgf000033_0001
あるこ とを特徴とする歪時効硬化特性に優れた深絞り用冷延鋼板である。 The cold-rolled steel sheet of the present invention has the above composition and a TSX r value ≥
Figure imgf000033_0001
This is a cold-rolled steel sheet for deep drawing, which is characterized by having excellent strain age hardening characteristics.
T S X r値が 750MPaを下回る鋼板では、 構造部材的な要素をもつ部材に広く適用 することができない。 また、 さらに適用範囲を拡げるには T S X r値は 850MPa以上 とするのが好ましい。  If the TSXr value is less than 750MPa, it cannot be widely applied to structural members. Further, in order to further expand the applicable range, the TSXr value is preferably set to 850 MPa or more.
従来の塗装焼付け処理条件は、 nO . ^ X SOiuin が標準として採用されている。 な お、 多量の固溶 Nを含む本発明鋼板に 5 %以上の歪が加わる場合は、 より緩やかな (低温側の） 処理でも硬化が達成され、 言い換えれば時効条件をより幅広く とるこ とが可能である。 また、 一般に、 硖化量を稼ぐには、 過度の時効で軟化させない限 りにおいて、 より高温で、 より長時間保持することが有利である。 As the conventional paint baking conditions, nO. ^ X SOiuin is adopted as a standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, hardening is achieved even with milder (lower temperature) treatment, in other words, aging conditions can be broadened. And it is possible. In general, in order to increase the amount of aging, it is advantageous to hold at a higher temperature and for a longer time as long as the material is not softened by excessive aging.
具体的に述べると、 本発明鋼板では、 予変形後に硬化が顕著となる加熱温度の下 限は概ね 100 である。 一方、 加熱温度が 300 °Cを超えると硬化が頭打ちとなり、 逆にやや軟化する傾向が現れるほか、 熱歪やテンパーカラーの発生が目立つように なる。 また、 保持時間については、 加熱温度 200 °C程度のとき概ね 30 s程度以上と すれば略十分な硬化が達成される。 さらに大きな安定した硬化を得るには保持時間 6 0 s以上とするのが好ましい。 しかし、 20min を超える保持では、 さらなる硬化を望 みえないばかりか、 生産効率も著しく低下して実用面では不利である。  Specifically, in the steel sheet of the present invention, the lower limit of the heating temperature at which hardening is remarkable after pre-deformation is approximately 100. On the other hand, when the heating temperature exceeds 300 ° C, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous. If the holding time is about 30 s or more when the heating temperature is about 200 ° C, almost sufficient curing can be achieved. In order to obtain even greater stable curing, the holding time is preferably 60 s or more. However, holding for more than 20 min not only does not allow further curing, but also significantly reduces production efficiency, which is disadvantageous in practical use.
以上のことから、 本発明では、 時効処理条件として従来の塗装焼付処理条件の加 熱温度である 170 °C、 保持時間を 20min で評価すると定めた。 従来の塗装焼付け型 鋼板では十分な硬化が達成されない低温加熱 ·短時間保持の時効処理条件下でも、 本発明鋼板では大きな硬化が安定的に達成される。 なお、 加熱の仕方はとくに制限 されず、 通常の塗装焼付けに採用されている炉による雰囲気加熱のほか、 たとえば 誘導加熱や、 無酸化炎、 レーザ、 プラズマなどによる加熱などのいずれも好ましく 用いうる。 また、 強度を上昇させたい部分のみを選択的に加熱してもよい。  From the above, in the present invention, it was determined that the aging treatment conditions were evaluated at a heating temperature of 170 ° C. and a holding time of 20 min under the conventional paint baking treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention. The method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used. Alternatively, only the portion where the strength is to be increased may be selectively heated.
自動車用の部品強度は外部からの複雑な応力負荷に抗しうる必要があり、 それゆ え素材鋼板では小さな歪域での強度特性だけでなく大きな歪域での強度特性も重要 となる。 本癸明者らはこの点に鑑み、 自動車部品の素材となすべき本発明鋼板の B Hを 80MPa 以上とするとともに、 A T Sを 40MPa 以上とする。 なお、 より好ましく は、 B H IOOMPa以上、 厶 T S 50MPa 以上とする。 B Hと A T Sをより大きくするに は、 時効処理の際の加熱温度をより高温側に、 および または、 保持時間をより長 時間側に、 設定すればよい。  The strength of automotive components needs to be able to withstand complex external stress loads, so that not only the strength characteristics in a small strain range but also the strength characteristics in a large strain range are important for a material steel plate. In view of this point, the present inventors set the BH of the steel sheet of the present invention, which is to be used as a material for automobile parts, to be 80 MPa or more and the ATS to be 40 MPa or more. More preferably, the pressure is BHIOOMPa or more, and the temperature TS50MPa or more. To increase BH and ATS, the heating temperature and / or the holding time during the aging treatment may be set to a higher temperature and / or a longer time.
また、 本発明鋼板は、 成形加工されない状態では、 室温で 1年程度の長時間放置 されても時効劣化 （Y Sが増加しかつ E 1が減少する現象） は起こらないという、 従来にない利点が備わっている。 また、 本発明では、 上記した本発明冷延鋼板の表面に溶融亜鉛めつきまたは合金 化溶融亜鉛めつきを施しても何ら問題はなく、 めっき前と同等位の TS、 BH、 厶 TSを示す。 また、 溶融亜鉛めつきおよび合金化溶融亜鉛めつき以外のめっきの種 類としては、 電気亜鉛めつき、 電気錫めつき、 電気クロムめつき、 電気ニッケルめ つき等、 いずれも好ましく適用しうる。 第 4の本発明による製造条件について述べる。 In addition, the steel sheet of the present invention has an unprecedented advantage that it does not undergo aging deterioration (phenomenon in which YS increases and E 1 decreases) even when it is left at room temperature for a long period of about one year when it is not formed. Equipped. Further, in the present invention, there is no problem even if hot dip galvanizing or alloyed hot dip galvanizing is performed on the surface of the above-mentioned cold rolled steel sheet of the present invention, and TS, BH, and mu TS are comparable to those before plating. . Further, as a plating type other than the hot-dip galvanizing and the alloyed hot-dip galvanizing, any of electroplating, electroplating, electrochromic plating, and electroplating can be preferably used. Fourth manufacturing conditions according to the present invention will be described.
まず、 C ： 0.01 %未満、 N： 0.0050 〜0.04%、 A1： 0.005 〜0.03 %、 Si： 0,005 ~1.0 %、 Mn : 0.01〜1.5 %、 P ： 0.1%以下、 S : 0.01%以下を含み、 あるいは さらに B ： 0.0001〜0.003%と共に、 Nb : 0.005 〜0.050 %、 Ti： 0.005 〜0.070 %、 V : 0.005 〜0.10%のうち 1種または 2種以上を含み、 かつ N/(A1+Nb+Ti+V+B) : 0.3 0以上になる組成を有する鋼が転炉等の通常公知の溶製法で溶製され、 造塊法あるい は連続鎳造法で凝固させられて鋼素材となる。 First, C: less than 0.01%, N: 0.0050 to 0.04%, A1: 0.005 to 0.03%, Si: 0.005 to 1.0%, Mn: 0.01 to 1.5%, P: 0.1% or less, S: 0.01% or less, Or B: 0.001 to 0.003%, Nb: 0.005 to 0.050%, Ti: 0.005 to 0.070%, V: 0.005 to 0.10%, and N / (A1 + Nb + Ti + V + B): Steel having a composition of not less than 0.30 is smelted by a commonly known smelting method such as a converter and solidified by an ingot-making method or a continuous sintering method to form a steel material.
この鋼素材は加熱、 均熱ののち熱間圧延されて熱延板となる。 加熱温度 （SRT) が低すぎると加工性の改善効果が飽和し、 しかも熱間圧延時の圧延負荷が増大し圧 延トラブルが発生したり、 また固溶 Nの均一化不足を招くおそれが生じてくるので、 SRTは 950 で以上が好ましい。 なお、 深絞り性向上のためには固溶 Cを固定し炭 化物として析出させておくのが有利であり、 それには S RTは 1300で以下が好まし い。 なお加工性をより一層の向上させるには 1150。C以下とするのが好ましい。  This steel material is heated and soaked, then hot-rolled into a hot-rolled sheet. If the heating temperature (SRT) is too low, the effect of improving workability saturates, and the rolling load during hot rolling increases, which may cause rolling troubles and may cause insufficient uniformity of solid solution N. The SRT should be 950 or higher. In order to improve the deep drawability, it is advantageous to fix solid solution C and precipitate it as a carbide. For this purpose, the SRT is 1300 and the following is preferable. To further improve workability, 1150. C or less is preferable.
熱間圧延の粗圧延〜仕上圧延の全圧下率は、 80%未満であると熱延板の結晶粒微 細化が不十分となるため、 80%以上とするのが好ましい。  If the total rolling reduction from the rough rolling to the finish rolling in the hot rolling is less than 80%, the grain refinement of the hot-rolled sheet becomes insufficient, so that it is preferably at least 80%.
粗圧延温度が lOOOt超であると y→α変態粒が粗大化して r値が低下し、 Ar3 未 満であると α粒が再結晶粗大化または粒成長することにより r値が低下するので、 粗圧延は 1000で以下 Ar3 以上の温度域で行うことが好ましい。 Rough rolling temperature to y → alpha transformation grains coarsened to be lOOOt than reduces the r value, the r value is decreased by alpha particles and Ar 3 is less than grows recrystallization coarsening or grain , rough rolling is preferably carried out in a temperature range of less than the Ar 3 in 1000.
一方、仕上圧延を Αι·3 超の温度域で終えると γ→α変態により集合組織がランダ ム化し、 優れた深絞り性が得られず、 一方、 仕上圧延を 600 °C未満で終えても、 よ り一層の深絞り性の向上は望めず圧延荷重が増大するのみとなるので、 仕上圧延は A r s 以下 600 以上の温度域で行うことが好ましい。 On the other hand, when finish rolling is completed in a temperature range exceeding Αι · 3, the texture becomes random due to γ → α transformation, and excellent deep drawability cannot be obtained.On the other hand, even when finish rolling is completed below 600 ° C , Yo Since further improvement in deep drawability cannot be expected and only the rolling load increases, the finish rolling is preferably performed in a temperature range of Ars to 600 or more.
また、 仕上圧延時に潤滑圧延を行わないと、 ロールと鋼板との間の摩擦力により、 鋼板表層部に付加的剪断力が働き、 その結果、 鋼板表層部に深絞り性に好ましくな い { 110 } 方位が優先的に形成されるために、 深絞り性が劣化する。 よって、 仕上 圧延は潤滑しつつ行うのが好ましい。  If lubrication rolling is not performed during finish rolling, additional shearing force acts on the surface of the steel sheet due to the frictional force between the roll and the steel sheet. } Since the orientation is preferentially formed, deep drawability deteriorates. Therefore, the finish rolling is preferably performed while lubricating.
ついで、 熱延板はコイル状に巻き取られる。 なお、 卷取工程を経た被処理材はコ ィルとも呼ばれる。 熱延板の卷取温度 （C T ) は、 高温ほど炭化物の粗大化には有 利であるが、 800 を超えると熱延板表面に形成されるスケールが厚くなりスケー ル除去作業の負荷が増大したり、 窒化物形成が進行してコイル長手方向の固溶 N量 の変動を招き、 一方、 400 未満では卷取作業が困難になる。 このため C Tは 800 〜400 tとするのが好ましい。  Next, the hot rolled sheet is wound into a coil. The material to be processed after the winding step is also called a coil. The winding temperature (CT) of the hot-rolled sheet is higher at higher temperatures, which is advantageous for coarsening of carbides. However, when it exceeds 800, the scale formed on the hot-rolled sheet surface becomes thicker and the load of scale removal work increases. Or the formation of nitrides progresses, causing a change in the amount of solute N in the longitudinal direction of the coil. On the other hand, if it is less than 400, winding becomes difficult. For this reason, CT is preferably set to 800 to 400 t.
次いで、 得られた熱延板を、 連続焼鈍またはパッチ焼鈍により再結晶焼鈍する。 この焼鈍 （熱延板焼鈍） は、 仕上圧延で行った α域温間圧延により形成された圧延 加工集合組織を再結晶させて、 再結晶集合組織を得るために行うものである。  Next, the obtained hot rolled sheet is subjected to recrystallization annealing by continuous annealing or patch annealing. This annealing (hot rolled sheet annealing) is performed in order to obtain a recrystallized texture by recrystallizing the rolled texture formed by the α region warm rolling performed in the finish rolling.
次いで、 熱延板は冷間圧延されて冷延板となる。 冷間圧延の圧下率は、 60%未満 では高 r値が期待できず、 一方、 95%を超えると r値がかえって低下するので、 60 ~95%とするのが好ましい。  Next, the hot-rolled sheet is cold-rolled to be a cold-rolled sheet. If the rolling reduction of the cold rolling is less than 60%, a high r-value cannot be expected, while if it exceeds 95%, the r-value rather decreases, so that it is preferable to be 60 to 95%.
次いで、 冷延板は再結晶焼鈍される。 この焼鈍は、 連続焼鈍ライン、 連続溶融亜 鉛めつきラインのいずれかで行うのが好ましい。 焼鈍条件は焼鈍温度 650 で以上 X 保持時間 5秒以上とするのが好ましい。 焼鈍温度 650 で以上、 保持時間 5秒以上の いずれかが満たされないと再結晶が完了せず深絞り性が低下する。 なお、 さらに優 れた深絞り性を得るには、 焼鈍温度 800 ^以上 X保持時間 5秒以上が好ましい。 た だし、 焼鈍温度が 900 を超えると炭化物の再溶解が進行し固溶 Cが過度に増加す るため遅時効性 （耐常温時効性） が低下し、 さらに、 α→γ変態が生じた場合は集 合組織がランダム化して r値が低下し深絞り性が損なわれるため、 焼鈍温度は 900 °C以下とするのが好まレぃ。 Next, the cold rolled sheet is subjected to recrystallization annealing. This annealing is preferably performed in either a continuous annealing line or a continuous molten zinc plating line. The annealing conditions are preferably annealing temperature of 650 or more and X holding time of 5 seconds or more. Unless any of the annealing temperature of 650 or more and the holding time of 5 seconds or more are satisfied, recrystallization is not completed and the deep drawability decreases. In order to obtain even better deep drawability, the annealing temperature is preferably 800 ^ or more and the X holding time is preferably 5 seconds or more. However, when the annealing temperature exceeds 900, the re-dissolution of carbides proceeds and the solute C increases excessively, so that the slow aging (normal temperature aging resistance) is reduced and the α → γ transformation occurs. Since the texture is randomized and the r-value is reduced and the deep drawability is impaired, the annealing temperature is 900 It is preferable that the temperature be below ° C.
さらに、 冷延鋼板を再結晶焼鈍して得られた冷延焼鈍板には、 必要に応じて溶融 亜鉛めつき、 あるいはさらに合金化処理を施すが、 その場合、 めっき処理では、 再 結晶焼鈍後からめっき処理前までの冷却速度を 5 /s以上とし、 溶融亜鉛めつきす る時の板温を 400〜600 °Cとするのが好ましく、 合金化処理では、 処理温度を 400 〜600 でとし、 処理時間を 5〜40秒とするのが好ましい。  Furthermore, the cold-rolled annealed sheet obtained by recrystallizing and annealing the cold-rolled steel sheet is subjected to hot-dip galvanizing or further alloying as necessary. It is preferable that the cooling rate from the time before to the plating process is 5 / s or more, and the sheet temperature when hot-dip galvanizing is 400 to 600 ° C. In the alloying process, the processing temperature is 400 to 600 ° C. The processing time is preferably 5 to 40 seconds.
なお、 再結晶焼鈍後の冷延鋼板あるいは溶融亜鉛めつき鋼板は、 形状矯正、 表面 粗さ調整のためにこれを調質圧延してもよい。 この調質圧延の圧下率は 10%以下が 好ましい。 この圧下率が 10%を超えると r値が低下するからである。 第 5の本発明の高張力冷延鋼板の組成限定理由について説明する。  The cold-rolled steel sheet or the hot-dip galvanized steel sheet after the recrystallization annealing may be subjected to temper rolling for shape correction and surface roughness adjustment. The rolling reduction of this temper rolling is preferably 10% or less. This is because when the rolling reduction exceeds 10%, the r value decreases. The reason for limiting the composition of the high-tensile cold-rolled steel sheet of the fifth invention will be described.
C： 0. 0015〜0. 025 %  C: 0.0015 to 0.025%
cは、 組織を均一かつ微細に制御し、 ァシキユラ一フェライ ト相を十分な量確保 するため、 本発明では 0. 0015%以上含有する必要がある。 一方、 0. 025 %を超える と、 鋼板中の炭化物分率が過大となり、 延性、 r値さらには成形性が顕著に低下す る。 このようなことから、 Cは 0. 0015~0. 025 %の範囲内に限定した。 なお、 成形 性の向上という観点からは、 0. 020 %以下とするのが好ましく、 より好ましくは 0. 0 10 %以下である。 また、 特に B H量および材質を安定性させる観点からは、 C含有 量は （12/ 93) Nb (%) 超え （ここで、 Nbは Nb含有量 （％) ) とするのがより好ま しい。  In the present invention, c must be contained at least 0.0015% in order to control the structure uniformly and finely and to secure a sufficient amount of the ferro-ferrite phase. On the other hand, when the content exceeds 0.025%, the carbide fraction in the steel sheet becomes excessive, and the ductility, the r value, and the formability are significantly reduced. For these reasons, C is limited to the range of 0.0015 to 0.025%. From the viewpoint of improving the moldability, the content is preferably at most 0.020%, more preferably at most 0.010%. Also, from the viewpoint of stabilizing the BH content and the material quality, it is more preferable that the C content exceeds (12/93) Nb (%) (where Nb is the Nb content (%)).
Si： 1. 0 %以下  Si: 1.0% or less
Siは、 鋼の延性を顕著に低下させることなく、 鋼板を高強度化させることができ る有用な元素であり本発明では 0. 005 %以上含有するのが好ましく、 特に高強度が 必要な場合は 0. 10%以上含有するのがより好ましい。 一方、 Siは、 熱間圧延時に変 態点を大きく上昇させて品質、 形状の確保を困難にしたり、 あるいはまた表面性状、 化成処理性など、 特に鋼板表面の美麗性に悪影響を与え、 さらにめつき性にも悪影 響を及ぼす元素であり、 本発明では 1. 0 %以下に限定した。 Siが 1. 0 %以下であれ ば、 上記した悪影響を低く抑えることができる。 なお、 特にめつき鋼板表面の美麗 性が要求される用途には、 Siは 0. 5 %以下とすることが望ましい。 Si is a useful element that can increase the strength of a steel sheet without significantly reducing the ductility of the steel, and is preferably contained in the present invention in an amount of 0.005% or more, particularly when high strength is required. Is more preferably 0.10% or more. On the other hand, Si significantly raises the transformation point during hot rolling, making it difficult to secure quality and shape, or adversely affects the surface properties and chemical conversion treatment, especially the aesthetics of the steel sheet surface. Poor influence on stickiness It is an element that has an effect, and is limited to 1.0% or less in the present invention. If the Si content is 1.0% or less, the above-mentioned adverse effects can be suppressed. It is desirable that the content of Si be 0.5% or less, particularly for applications requiring beautiful surface of the coated steel sheet surface.
Mn： 2. 0 %以下  Mn: 2.0% or less
Mnは、 Sによる熱間割れを防止する有効な元素であり、 含有する S量に応じて添 加するのが好ましく、 また Mnは結晶粒の微細化に対し大きな効果があり、 添加して 材質改善に利用することが望ましい。 Sを安定して固定するという観点から、 Mnは 0. 1 %以上の含有が望ましい。 また Mnは鋼板強度を増加させる元素であり、 より強度 が要求される場合には 0. 5 %以上含有するのが望ましい。 なお、 より好ましくは 0. 8 %以上である。  Mn is an effective element for preventing hot cracking due to S. It is preferable to add Mn in accordance with the amount of S contained.Mn has a great effect on refining crystal grains, and Mn is added. It is desirable to use it for improvement. From the viewpoint of stably fixing S, the content of Mn is desirably 0.1% or more. Mn is an element that increases the strength of the steel sheet, and when more strength is required, it is desirable to contain 0.5% or more. It is more preferably at least 0.8%.
Mn含有量をこのレベルまで高めると、 熱延条件の変動に封する鋼板の機械的性質、 とくに歪時効硬化特性のばらつきが顕著に改善されるという大きな利点がある。 し かし、 Mnを 2. 0 %を超えて過度に含有すると、 詳細な機構は不明であるが熱間変形 抵抗を增加させる傾向があり、 また溶接性、 溶接部成形性を劣化させる傾向となり、 さらにはフェライ トの生成が顕著に抑制され延性が顕著に低下し、 また r値が低減 する傾向が顕著となるため、 Mnは 2. 0 %以下に限定した。 より良好な耐食性と成形 性が要求される用途では、 1. 5 %以下とするのが好ましい。  When the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet to be sealed against the fluctuation of the hot rolling conditions, especially the variation of the strain aging hardening characteristics are remarkably improved. However, if Mn is excessively contained in excess of 2.0%, the detailed mechanism is unknown, but it tends to increase the hot deformation resistance, and also tends to deteriorate the weldability and the weldability. Furthermore, Mn was limited to 2.0% or less because the formation of ferrite was remarkably suppressed, ductility was remarkably reduced, and the r-value tended to be remarkably reduced. For applications requiring better corrosion resistance and formability, the content is preferably 1.5% or less.
P ： 0. 1%以下  P: 0.1% or less
Pは、 鋼の固溶強化元素として有用な元素であり、 強度増加の観点から 0. 002 %以上含有するのが好ましく、 特に高強度が必要な場合はより好ましくは 0. 02 %以上含有するのがより好ましい。 一方、 過度に含有すると、 鋼を脆化させ、 さら に鋼板の伸ぴフランジ加工性を悪化させる。 また、 Pは鋼中で偏析する傾向が強い ためそれに起因した溶接部の脆化をもたらす。 このため、 Pは 0. 1%以下に限定した。 なお、 伸ぴフランジ加工性や溶接部靱性がとくに重要視される用途では Pは 0. 08% 以下とするのが好ましい。 より好ましくは 0. 06%以下である。  P is a useful element as a solid solution strengthening element for steel, and is preferably contained in an amount of 0.002% or more from the viewpoint of increasing strength, and more preferably contained in an amount of 0.02% or more when high strength is required. Is more preferred. On the other hand, if it is contained excessively, it makes the steel brittle and further deteriorates the stretch flangeability of the steel sheet. In addition, P has a strong tendency to segregate in steel, resulting in embrittlement of the weld. Therefore, P was limited to 0.1% or less. In applications where stretch flangeability and weld toughness are particularly important, P is preferably set to 0.08% or less. More preferably, it is not more than 0.06%.
S ： 0. 02%以下 Sは、 鋼板中では介在物として存在し、 鋼板の延性を減少させ、 さらには耐食性 の劣化をもたらす元素であり、できるだけ低減するのが好ましく、本発明では Sは 0. 02%以下に限定した。 とくに、 良好な加工性が要求される用途には、 Sは 0. 015 % 以下とすることが好ましい。 また、 とくに優れた伸びフランジ加工性が要求される 場合には、 Sは 0. 010 %以下とすることが好ましい。 また、 詳細な機構は不明であ るが、 鋼板の歪時効硬化特性を安定して高いレベルに維持するためには、 Sを 0. 008 %以下まで低減するのが有効である。 S: 0.02% or less S is an element that exists as an inclusion in the steel sheet and reduces the ductility of the steel sheet, and furthermore, deteriorates the corrosion resistance.It is preferable that S is reduced as much as possible.In the present invention, S is limited to 0.02% or less. . In particular, for applications requiring good workability, S is preferably set to 0.015% or less. When particularly excellent stretch flangeability is required, S is preferably set to 0.010% or less. Although the detailed mechanism is unknown, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening property of the steel sheet at a high level.
A1： 0. 02%以下  A1: 0.02% or less
A1は、 脱酸剤として作用し鋼の清浄度を向上させ、 さらには鋼板の組織を微細化 する元素であり、 本発明では 0. 001 %以上の含有が望ましい。 本発明においては、 固溶状態の Nを強化元素として利用するが、 適正範囲の A1を含有したアルミキルド 鋼のほうが、 A1を添加しない従来のリムド鋼に比して、 機械的性質が優れている。 —方、 過剰の A1含有は、 鋼板の表面性状を惠化させ、 さらに固溶状態の Nを顕著に 低下させて、 本発明が主眼とする極めて大きな歪時効硬化量を得ることが困難とな る。 このようなことから、 本発明では A1は 0, 02%以下に限定した。 なお、 材質の安 定性という観点からは、 A1は 0. 001 〜0. 015 %とするのがさらに望ましい。 また、 A 1含有量の低減は結晶粒の粗大化につながる懸念もあるが、本発明では他の合金元素 を最適量とすることと、 焼鈍条件を最適な範囲とすることにより、 これを有効に防 止している。  A1 is an element that acts as a deoxidizing agent, improves the cleanliness of steel, and further refines the structure of the steel sheet. In the present invention, the content of A1 is preferably 0.001% or more. In the present invention, N in a solid solution state is used as a strengthening element.Aluminum-killed steel containing A1 within an appropriate range has better mechanical properties than conventional rimmed steel without A1 added. . On the other hand, an excessive amount of A1 enhances the surface properties of the steel sheet and further significantly reduces the N in the solid solution state, making it difficult to obtain an extremely large amount of strain age hardening which is the main objective of the present invention. You. For this reason, in the present invention, A1 is limited to 0.02% or less. From the viewpoint of the stability of the material, it is more desirable that A1 is 0.001 to 0.015%. In addition, although there is a concern that the reduction of the A1 content may lead to coarsening of crystal grains, the present invention makes it possible to optimize this by optimizing the amount of other alloying elements and setting the annealing conditions in the optimum range. Is prevented.
N ： 0. 0050〜0. 0250%  N: 0.0050 to 0.0250%
Nは、 固溶強化と歪時効硬化により鋼板の強度を増加させる元素であり、 本発明 において最も重要な元素である。 また、 本発明では、 適量の Nを含有して、 さらに 上記したように A1含有量を適正値に調整し、 さらには熱延条件、 焼鈍条件などの製 造条件を制御することにより、 冷延製品あるいはめっき製品で必要かつ十分な固溶 状態の Nを確保する。 これにより、 固溶強化と歪時効硬化による強度 （降伏応力お よび引張り強さ） 上昇効果が十分に発揮され、 引張強さ 340MPa以上、 焼付け硬化量 ( B H量） 80MPa 以上、 歪時効処理前後での引張強さの増加量 Δ T S 40MPa 以上、 という本発明鋼板の機械的性質の目標値を安定して得ることができる。 また、 Nは 変態点を低下させる作用を有し、 変態点を大きく割り込む圧延をしたくない薄物の 圧延等の場合に含有すると有効である。 N is an element that increases the strength of the steel sheet by solid solution strengthening and strain age hardening, and is the most important element in the present invention. Further, in the present invention, the cold rolling is performed by containing an appropriate amount of N, adjusting the A1 content to an appropriate value as described above, and further controlling the manufacturing conditions such as hot rolling conditions and annealing conditions. Ensure N in the solid solution state necessary and sufficient for the product or plating product. As a result, the effect of increasing the strength (yield stress and tensile strength) due to solid solution strengthening and strain age hardening is fully exhibited, with a tensile strength of 340 MPa or more and baking hardening amount. It is possible to stably obtain a target value of the mechanical properties of the steel sheet of the present invention, that is, (BH amount) 80 MPa or more, and an increase in tensile strength ΔTS 40 MPa or more before and after the strain aging treatment. In addition, N has an effect of lowering the transformation point, and it is effective to include N in the case of rolling of a thin material or the like in which rolling that does not want to greatly reduce the transformation point is desired.
Nが 0. 0050%未満では、 上記した強度上昇効果が安定して現れにくい。 一方、 N が 0. 0250%を超えると、 鋼板の内部欠陥発生率が高くなるとともに、 連続铸造時の スラブ割れなどの多発するようになる。 このため、 Nは 0. 0050〜0. 0250%の範囲に 限定した。 なお、 製造工程全体を考慮した材質の安定性，歩留り向上という観点か らは、 Nは 0. 0070〜0. 0200%、 さらに好ましくは 0. 0100〜0. 0170%の範囲とするの が好ましい。 なお、 本発明の範囲内の N量であれば、 溶接性等への悪影響は全くな い。  If N is less than 0.0050%, the above-described effect of increasing the strength is unlikely to appear stably. On the other hand, if N exceeds 0.0250%, the occurrence rate of internal defects in the steel sheet increases, and slab cracks and the like during continuous production become more frequent. For this reason, N was limited to the range of 0.0050 to 0.0250%. From the viewpoints of material stability and yield improvement taking into account the entire manufacturing process, N is preferably in the range of 0.0070 to 0.0200%, more preferably 0.0100 to 0.0170%. . If the N content is within the range of the present invention, there is no adverse effect on weldability and the like.
固溶状態の N： 0. 0010%以上  N in solid solution state: 0.0010% or more
冷延製品で十分な強度が確保され、 さらに Nによる歪時効硬化が有効に発揮され るには、 鋼板中に固溶状態の N (固溶 Nともいう） が少なくとも 0. 0010%以上存在 する必要がある。  In order for a cold-rolled product to have sufficient strength and for strain age hardening to be effectively exhibited by N, at least 0.0010% or more of N in solid solution (also called solid solution N) exists in the steel sheet There is a need.
ここで、 固溶 N量は、 鋼中の全 N量から、 析出 N量を差し引いた値を固溶 N量と する。 なお、 析出 N量の分析法としては、 本発明者らが種々の方法を比較検討した 結果、 定電位電解法を用いた電解抽出分析法により求めるのが有効である。 なお、 抽出分析に用いる地鉄を溶解する方法として、 酸分解法、 ハロゲン法および電解法 がある。 この中で、 電解法は炭化物、 窒化物などの極めて不安定な析出物を分解す ることなく、 安定して地鉄のみを溶解できる。 電解液としては、 ァセチル 'ァセト ン系を用いて、 定電位にて電解する。 本発明では定電位電解法を用いて析出 N量を 測定した結果が、 実際の材質の変化とよい対応を示した。  Here, the amount of solute N is defined as a value obtained by subtracting the amount of precipitated N from the total amount of N in steel. As a method of analyzing the amount of precipitated N, it is effective that the method is determined by an electrolytic extraction analysis method using a constant potential electrolysis method as a result of comparative studies of various methods by the present inventors. In addition, there are acid decomposition method, halogen method, and electrolysis method as a method of dissolving ground iron used for extraction analysis. Among these, the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides. Electrolyte at a constant potential using an acetyl-acetonic electrolyte. In the present invention, the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed a good correspondence with the actual material change.
このようなことから、 本発明では、 定電位電解法により抽出した残渣を化学分析 して残渣中の N量を求め、 これを析出 N量とする。  For this reason, in the present invention, the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N.
なお、 さらに高い B H量、 A T Sが必要な場合には、 固溶 N量を 0. 0020%以上、 さらに高い値を得るためには、 0.0030%以上とするのが好ましい。 固溶 N量の上限 値は特に限定しないが、 全 N量がすべて残留しても機械的性質の低下は小さい。 If a higher BH or ATS is required, the solid solution N content should be 0.0020% or more, In order to obtain a higher value, the content is preferably set to 0.0030% or more. The upper limit of the amount of solute N is not particularly limited, but even if all the amount of N remains, the decrease in mechanical properties is small.
NZAl (N含有量と A1含有量の比） ： 0.3 以上  NZAl (N content and A1 content ratio): 0.3 or more
製品状態で、 固溶 Nを 0.0010%以上安定させて残留させるためには、 Nを強力に 固定する元素である A1の量を制限する必要がある。 本発明の組成範囲内の N含有量 (0.0050〜0.0250%) と A1含有量 （0.02%以下） の組合せを広範囲に変えた鋼板に ついて検討した結果、 を 0.3 以上とすることにより、 冷延製品おょぴめっき 製品での固溶 Nを安定して 0.0010%以上とすることができることがわかった。 この ため、 を 0.3 以上に限定した。 なお、 歪時効硖化特性を安定して高めるとい う観点からは NZA1は 0.6 以上とするのが好ましい。 さらに好ましくは 0.8 以上で ある。  In the product state, it is necessary to limit the amount of A1, which is an element that strongly fixes N, in order to stably retain solute N of 0.0010% or more in the product state. As a result of examining a steel sheet in which the combination of the N content (0.0050 to 0.0250%) and the A1 content (0.02% or less) within the composition range of the present invention was changed over a wide range, by setting to 0.3 or more, the cold rolled product It was found that the solid solution N in the bath plating product can be stably increased to 0.0010% or more. Therefore, was limited to 0.3 or more. NZA1 is preferably set to 0.6 or more from the viewpoint of stably improving the strain aging characteristics. More preferably, it is 0.8 or more.
Nb： 0.002〜0.050 %  Nb: 0.002 to 0.050%
Nbは、 Bと複合してァシキユラ一フェライ ト相を生成することに有効に作用すし、 本発明では 0.002 %以上の含有を必要とする。 一方、 0.050 %を超えて含有すると、 効果が飽和するうえ、 熱間変形抵抗が顕著に増加し、 熱間圧延が困難となる。 この ため、 Nbは 0.002 -0.050 %の範囲内に限定した。 なお、 より好ましくは、 0.005 〜0.040 %である。  Nb effectively acts to form an ash-ferrite phase in combination with B, and the present invention requires a content of 0.002% or more in the present invention. On the other hand, if the content exceeds 0.050%, the effect is saturated and the hot deformation resistance is significantly increased, so that hot rolling becomes difficult. For this reason, Nb was limited to the range of 0.002-0.050%. In addition, more preferably, it is 0.005 to 0.040%.
B： 0.0001-0.0050%,  B: 0.0001-0.0050%,
Bは、 Nbと複合して、 ァシキユラ一フェライト相を生成することに有効に作用す る元素であり、 本発明では 0.0001%以上の含有を必要とする。 一方、 0.0050%を超 えて含有すると、歪時効硬化特性に寄与する固溶 Nを低減させる。 このため、 Bは 0· 0001~0.0050%の範囲内に限定した。 なお、 好ましくは、 0.0003〜0.0030%である。 より好ましくは、 0.0005〜0.0030%である。  B is an element that effectively acts to form the ferrite phase in combination with Nb. In the present invention, the content of 0.0001% or more is required. On the other hand, when the content exceeds 0.0050%, the solute N which contributes to the strain age hardening characteristics is reduced. Therefore, B is limited to the range of 0.0001% to 0.0050%. In addition, Preferably, it is 0.0003-0.0030%. More preferably, it is 0.0005 to 0.0030%.
本発明では、 上記した組成に加えてさらに、 次 a群〜 c群  In the present invention, in addition to the above composition, the following groups a to c
a群： Cu、 Ni、 Cr、 Moの 1種または 2種以上を合計で 1.0 %以下  Group a: 1.0% or less in total of one or more of Cu, Ni, Cr and Mo
b群： Ti、 Vの 1種または 2種以上を合計で 0.1 %以下 c群： Ca、 REM の 1種または 2種を合計で 0. 0010~0. 010 % Group b: 0.1% or less in total of one or more of Ti and V Group c: One or two of Ca and REM in total 0.0010 to 0.010%
の 1群または 2群以上を含有するのが好ましい。 It is preferable that one or two or more groups are contained.
a群の元素： Cu、 Ni、 Cr、 Moは、 いずれも鋼板の強度上昇に寄与する元素であり、 必要に応じ選択して単独または複合して含有できる。 このような効果は、 Cu、 Ni、 C r、 Moをそれぞれ 0. 01%以上の含有で認められる。 しかし、 含有量が多すぎると熱間 変形抵抗が増加し、 あるいは化成処理性や広義の表面処理特性が悪化するうえ、 溶 接部が硬化し溶接部成形性が劣化する。 このため、 Cu、 Ni、 Cr、 Moはそれぞれ単独 では 1. 0 %以下、 1. 0 %以下、 0. 5 %以下、 0. 2 %以下とするのが好ましく、 複合 して含有する場合には合計で 1. 0 %以下とするのが好ましい。  Group a elements: Cu, Ni, Cr, and Mo are all elements that contribute to the increase in the strength of the steel sheet, and can be selected singly or in combination as necessary. Such an effect is recognized when Cu, Ni, Cr, and Mo are contained at 0.01% or more, respectively. However, if the content is too large, the hot deformation resistance increases, or the chemical conversion property and the surface treatment properties in a broad sense deteriorate, and the welded part is hardened and the formability of the welded part is deteriorated. Therefore, it is preferable that Cu, Ni, Cr, and Mo each alone be 1.0% or less, 1.0% or less, 0.5% or less, and 0.2% or less. Is preferably 1.0% or less in total.
b群の元素： Ti、 Vは、 いずれも結晶粒の微細化 ·均一化に寄与する元素であり、 必要に応じ選択して単独または複合して含有できる。 このような効果は、 Ti、 Vを それぞれ 0. 005 %以上の含有で認められる。 しかし、 含有量が多すぎると、 熱間変 形抵抗が増加し、 また化成処理性や広義の表面処理特性が悪化する。 さらに固溶 N を低減する悪影響もある。 このため、 Ti、 Vは単独ではそれぞれ 0. 1%以下、 0. 1 % 以下とするのが好ましく、 複合して含有する場合には合計で 0. 1%以下とするのが好 ましい。  Group b elements: Both Ti and V are elements that contribute to the refinement and uniformization of crystal grains, and can be selected as necessary and contained alone or in combination. Such an effect is recognized when the contents of Ti and V are each 0.005% or more. However, if the content is too large, the hot deformation resistance increases, and the chemical conversion property and the surface treatment properties in a broad sense deteriorate. In addition, there is an adverse effect of reducing solid solution N. Therefore, Ti and V alone are preferably 0.1% or less and 0.1% or less, respectively, and when they are contained in combination, the total content is preferably 0.1% or less.
c群の元素： Ca、 REM は、 いずれも介在物の形態制御に役立つ元素であり、 特に 伸ぴフランジ成形性の要求がある場合には、 単独または複合して含有するのが好ま しい。 d群の元素の合計で、 0. 0010%未満では介在物の形態制御効果が不足し、 一 方、 0. 010 %を超えると表面欠陥の発生が目立つようになる。 このため、 d群の元 素を合計で 0. 0010〜0. 010 %の範囲に限定することが好ましく、 これにより、 表面 欠陥の発生を伴うことなく伸びフランジ加工性を改善することができる。  Group c elements: Ca and REM are both elements that are useful for controlling the morphology of inclusions, and if there is a requirement for stretch-flange formability, they are preferably contained alone or in combination. If the total of the elements in group d is less than 0.0010%, the effect of controlling the morphology of inclusions will be insufficient, whereas if it exceeds 0.010%, the occurrence of surface defects will be noticeable. For this reason, it is preferable to limit the total of the elements of the d group to the range of 0.0010 to 0.010%, whereby the stretch flangeability can be improved without the occurrence of surface defects.
本発明鋼板の組織について説明する。  The structure of the steel sheet of the present invention will be described.
本発明鋼板は、 面積率で 5 %以上のァシキユラ一フェライ ト相と平均結晶粒径 20 /z m以下のフェライ ト相から成る組織を有する。  The steel sheet of the present invention has a structure composed of a ferrite phase having an area ratio of 5% or more and a ferrite phase having an average crystal grain size of 20 / zm or less.
ァシキユラ一フェライ ト相の面積率： 5 %以上 本発明の冷延鋼板は、 ァシキユラ一フェライ ト相を面積率で 5 %以上含有する。 ァシキユラ一フェライ ト相が 5 %以上存在することにより、 良好な延性と、 さらに、 大きな歪時効硬化量が得られる。 詳細な機構は不明であるが、 ァシキユラ一フェラ ィト相の存在により、 時効前の予歪み加工時に極めて有効に歪が内部に蓄積される ためと推定される。 さらに、 ァシキユラ一フェライ ト相の存在は、 常温での時効劣 化を改善し、 常温非時効性とするためにも有効である。 なお、 良好な強度一延性パ ランス、 より高い強度を得るためには、 ァシキユラ一フェライ ト相の面積率を 10% 以上とするのが好ましい。 なお、 20%を超える多量のァシキユラ一フェライ ト相の 存在は r値の低下という問題がある。 このため、 ァシキユラ一フェライ ト相の面積 率は、 5 %以上、 好ましくは 10%以上、 20%以下である。 Area ratio of ferrite phase: 5% or more The cold-rolled steel sheet of the present invention contains 5% or more in area ratio of the ferrite-ferrite phase. The presence of at least 5% of the ferrite phase provides good ductility and a large amount of strain age hardening. Although the detailed mechanism is unknown, it is presumed that the presence of the asymmetric ferrite phase causes the strain to be accumulated very effectively inside the pre-strain machining before aging. Furthermore, the presence of the ferrite phase is effective in improving the deterioration of aging at room temperature and making it non-aging at room temperature. In order to obtain a good strength-ductility balance and higher strength, it is preferable that the area ratio of the ferrite-ferrite phase is 10% or more. In addition, the presence of a large amount of ash-ferrite phase exceeding 20% has a problem that the r-value decreases. For this reason, the area ratio of the fermentation phase is 5% or more, preferably 10% or more and 20% or less.
本発明でいう、 ァシキユラ一フェライ ト相は、 本発明の組成のような極低炭素鋼 に特有な内部に炭化物を伴わない低温変態相で、 主として光学顕微鏡観察により通 常のポリゴナルフェライ トとは明確に識別可能であり、 内部の転位密度が高くポリ ゴナルフェライ ト相より硬質な相である。  The ferrite phase referred to in the present invention is a low-temperature transformation phase unique to ultra-low carbon steel without a carbide therein, such as the composition of the present invention, which is mainly a normal ferrite phase obtained by observation with an optical microscope. Is a phase that is clearly identifiable, has a high internal dislocation density, and is harder than the polygonal ferrite phase.
光学顕微鏡観察によれば、 ァシキユラ一フェライ ト相は、 ①粒界が不規則に角張 つた結晶粒状、 ②析出物のような粒界に添って存在する結晶粒状、 ③引つかき傷状 の模様を呈する結晶粒状または結晶粒群状 （比較的大きい第 2相粒中に亜粒界が多 数見られる） などのいずれかかが単独または複合して分布するもので、 これらは、 通常のポリゴナルフェライ トとは明確に区別できる。 また、 さらに粒内の腐食され た色調が、 マルテンサイ トやべイナイ トとは異なり、 通常のポリゴナルフェライ ト とはほとんど変わりないことから、 マルテンサイ トやべイナィ トとも明確に区別で きる。 透過型電子顕微鏡による観察によれば、 ァシキユラ一フェライ ト相は、 粒界 近傍および Zまたは粒内の転位密度が非常に高く、 とくに③の形態のものは転位密 度が非常に高い部分と比較的低い部分とが層'状となっている。  According to the observation with an optical microscope, the ferrite phase has the following features: (1) crystal grains in which the grain boundaries are irregularly angular, (2) crystal grains existing along the grain boundaries such as precipitates, and (3) scratch-like patterns. One or a combination of grains (a number of sub-boundaries are found in relatively large second-phase grains). It can be clearly distinguished from Nalferite. In addition, the corroded color tone in the grains is different from martensite and bainite, and is almost the same as ordinary polygonal ferrite, so that it can be clearly distinguished from martensite and bainite. According to the observation with a transmission electron microscope, the dispersoid density in the ferrite phase is very high near the grain boundary and in the Z or intragranular phase. The lower part is in the form of a layer.
本発明の冷延鋼板は、 高い成形性が要求される自動車用鋼板を対象としており、 延性を確保するために、 ァシキユラ一フェライト相以外の相はフェライ ト相とする。 フェライ ト相の面積率が 80%未満では、 加工性が要求される自動車用鋼板として必 要な延性と、 高い r値を確保することが困難となる。 なお、 さらに良好な延性が要 求される場合は、 フェライ ト相の面積率は 80%以上、 さらに好ましくは 85%以上と するのが望ましい。 なお、 本発明でいうフェライ トは、 歪みが残留していない状態 のいわゆるポリゴナルなフェライトをいうものとする。 The cold-rolled steel sheet of the present invention is intended for a steel sheet for automobiles that requires high formability, and in order to ensure ductility, a phase other than the ferrite phase is a ferrite phase. If the area ratio of the ferrite phase is less than 80%, it is difficult to secure the ductility necessary for an automotive steel sheet requiring workability and a high r-value. If better ductility is required, the area ratio of the ferrite phase is preferably at least 80%, more preferably at least 85%. The ferrite in the present invention refers to a so-called polygonal ferrite in which no distortion remains.
フェライト相の平均結晶粒径： 20 /i in 以下  Average grain size of ferrite phase: 20 / i in or less
本発明では平均結晶粒径として、 断面組織写真から A S T Mに規定された求積法 により算出した値と、同じく A S T Mに規定された切断法により求めた公称粒径（例 えば梅本ら：熱処理， 24 (1984) ，334参照） のうち、 より大きい方を採用する。 本発明の冷延鋼板では、 製品段階で所定量の固溶 N量を確保しているが、 本発明 者らの実験 ·検討によれば、 同一量の固溶 Nを有する鋼板でも、 歪時効硬化特性に ばらつきが生じる場合があり、 その主たる要因の一つが結晶粒径であることが判明 した。 本発明のような組織では、 平均結晶粒径を少なく とも 以下、 望ましく は 15 /z m 以下にすることにより、 安定して高い B H量、 A T Sが得られる。 詳細な 機構は不明であるが、 結晶粒界への合金元素の偏析と析出、 さらにはこれらに及ぼ す加工、 熱履歴の影響に関係していると推定される。  In the present invention, as the average crystal grain size, a value calculated from the cross-sectional structure photograph by the quadrature method specified by ASTM and a nominal grain size determined by the cutting method also specified by ASTM (for example, Umemoto et al. (1984), 334). In the cold-rolled steel sheet of the present invention, a predetermined amount of solute N is ensured at the product stage. However, according to experiments and studies conducted by the present inventors, even in a steel sheet having the same amount of solute N, strain aging has occurred. In some cases, the curing characteristics varied, and it was found that one of the main factors was the crystal grain size. In the structure as in the present invention, a high BH content and ATS can be obtained stably by setting the average crystal grain size to at least not more than, preferably not more than 15 / zm. Although the detailed mechanism is unknown, it is presumed to be related to the segregation and precipitation of alloying elements at the grain boundaries, as well as the effects of processing and thermal history on these.
したがって、 歪時効硬化特性の安定化を図るためには、 フェライト相の平均結晶 粒径を 20 /z m 以下、 好ましくは 15 # ιη 以下とすることが好ましい。  Therefore, in order to stabilize the strain age hardening characteristics, the average crystal grain size of the ferrite phase is preferably 20 / zm or less, and more preferably 15 # ιη or less.
上記した組成と組織を有する本発明の冷延鋼板は、 引張強さ （T S ) 340 MPa 以上で概ね 590MPa以下を有し、 さらに r値が 1. 2 以上の高 r値と、 優れた歪時効硬 化特性を有する冷延鋼板である。 T Sが 340MPaを下回る鋼板では、 構造部材的な要 素をもつ部材に広く適用することができない。 また、 さらに適用範囲を拡げるには T Sは 400MPa以上とするのが望ましい。 また、 r値が 1. 2 未満では、 広範囲なプレ ス成形部品に適用できない。 なお、 r値の好ましい範囲は 1. 3 以上である。  The cold-rolled steel sheet of the present invention having the above-described composition and structure has a tensile strength (TS) of 340 MPa or more and about 590 MPa or less, a high r value of r value of 1.2 or more, and excellent strain aging. It is a cold rolled steel sheet having hardening characteristics. Steel sheets with TS below 340MPa cannot be widely applied to structural members. In order to further expand the applicable range, it is desirable that T S be 400 MPa or more. If the r-value is less than 1.2, it cannot be applied to a wide range of press-formed parts. The preferred range of the r value is 1.3 or more.
従来の塗装焼付け処理条件は、 170 °C X 20min が標準として採用されている。 な お、 多量の固溶 Nを含む本発明鋼板に 5 %以上の歪が加わる場合は、 より緩やかな (低温側の） 処理でも硬化が達成され、 言い換えれば時効条件をより幅広く とるこ とが可能である。 また、 一般に、 硬化量を稼ぐには、 過度の時効で軟化させない限 りにおいて、 より高温で、 より長時間保持することが有利である。 The conventional paint baking condition is 170 ° C X 20 min as standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, Hardening is also achieved with (lower temperature) treatment, in other words, a wider range of aging conditions is possible. In general, in order to increase the amount of hardening, it is advantageous to hold at a higher temperature and for a longer time, unless softening is caused by excessive aging.
具体的に述べると、 本発明鋼板では、 予変形後に硬化が顕著となる加熱温度の下 限は概ね 100 °Cである。 一方、 加熱温度が 300 を超えると硬化が頭打ちとなり、 逆にやや軟化する傾向が現れるほか、 熱歪やテンパーカラーの発生が目立つように なる。 また、 保持時間については、 加熱温度 200 °C程度のとき概ね 30 s程度以上と すれば略十分な硬化が達成される。 さらに大きな安定した硕化を得るには保持時間 6 0 s以上とするのが好ましい。 しかし、 20min を超える保持では、 さらなる硬化を望 みえないばかり力 、 生産効率も著しく低下して実用面では不利である。  Specifically, in the steel sheet of the present invention, the lower limit of the heating temperature at which hardening becomes significant after pre-deformation is approximately 100 ° C. On the other hand, when the heating temperature exceeds 300, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous. If the holding time is about 30 s or more when the heating temperature is about 200 ° C, almost sufficient curing can be achieved. In order to obtain a still more stable deterioration, the holding time is preferably 60 s or more. However, holding for more than 20 min is not practical because it does not allow further curing and the production efficiency drops significantly.
以上のことから、 本発明では、 時効処理条件として従来の塗装焼付処理条件の加 熱温度である 170 °C、 保持時間を 20min で評価すると定めた。 従来の塗装焼付け型 鋼板では十分な硬化が達成されない低温加熱 ·短時間保持の時効処理条件下でも、 本発明鋼板では大きな硬化が安定的に達成される。 なお、 加熱の仕方はとくに制限 されず、 通常の塗装焼付けに採用されている炉による雰囲気加熱のほか、 たとえば 誘導加熱や、 無酸化炎、 レーザ、 プラズマなどによる加熱などのいずれも好ましく 用いうる。  From the above, in the present invention, it was determined that the aging treatment conditions were evaluated at a heating temperature of 170 ° C. and a holding time of 20 min under the conventional paint baking treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention. The method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used.
自動車用の部品強度は外部からの複雑な応力負荷に抗しうる必要があり、 それゆ え素材鋼板では小さな歪域での強度特性だけでなくより大きな歪域での強度特性も 重要となる。 本発明者らはこの点に鑑み、 自動車部品の素材となすべき本発明鋼板 の B H量 （比較的小さな歪域の強度特性に対応） を 80MPa 以上とするとともに、 Δ T S量 （比較的大きな歪域の強度特性に対応） を 40MPa 以上とする。 なお、 より好 ましくは、 B H量 lOOMPa以上、 Δ T S 50MPa 以上とする。 また、 時効処理の際の加 熱温度をより高温側に、 および Zまたは、 保持時間をより長時間側に、 設定するこ とにより、 B H量と A T Sをより大きくできる。  The strength of automotive components needs to be able to withstand complex external stress loads. Therefore, not only the strength characteristics in a small strain range but also the strength characteristics in a larger strain range are important for a material steel plate. In view of this point, the present inventors set the BH amount (corresponding to the strength characteristic in a relatively small strain range) of the steel sheet of the present invention to be used as a material for automobile parts to 80 MPa or more, and set the ΔTS (Corresponding to the strength characteristics of the region) should be 40MPa or more. It is more preferable that the BH content be lOOMPa or more, and ΔTS 50MPa or more. Also, by setting the heating temperature during aging treatment to a higher temperature side and setting Z or the holding time to a longer time side, the BH amount and ATS can be further increased.
ところで、 本発明の効果は製品板厚が比較的厚い場合でも癸揮されうるが、 製品 板厚が 3. 2mm を超える場合には、 冷延板焼鈍工程で必要十分な冷却速度を確保する ことができず、 連続焼鈍時に歪時効が生じ、 製品として目標とする歪時効硬化特性 が得にくくなる。 したがって、 本発明鋼板の板厚は 3. 2 mm以下とするのが好ましい。 また、 本発明では、 上記した本発明冷延鋼板の表面に電気めつきまたは溶融めつ きを施しても何ら問題はない。 これらめつき鋼板も、 めっき前と同程度の T S、 B H量、 A T S量を示す。 めっきの種類としては、 電気亜鉛めつき、 溶融亜鉛めつき、 合金化溶融亜鉛めつき、 電気錫めつき、 電気クロムめつき、 電気ニッケルめっき等、 いずれも好ましく適用しうる。 第 6の本発明による鋼板の製造方法について説明する。 By the way, the effect of the present invention can be obtained even when the product thickness is relatively large, If the sheet thickness exceeds 3.2 mm, it is not possible to secure the necessary and sufficient cooling rate in the cold-rolled sheet annealing process, and strain aging will occur during continuous annealing, and the strain aging hardening characteristics targeted as a product will be obtained. It becomes difficult. Therefore, the steel sheet of the present invention preferably has a thickness of 3.2 mm or less. Further, in the present invention, there is no problem even if the surface of the above-mentioned cold rolled steel sheet of the present invention is electroplated or melted. These plated steel sheets also show the same amount of TS, BH and ATS as before plating. As the type of plating, any of electroplating, hot-dip galvanizing, alloyed hot-dip galvanizing, electro-tin plating, electro-chrome plating, and electro-nickel plating can be preferably applied. A method for manufacturing a steel sheet according to the sixth invention will be described.
本発明鋼板は、 基本的には、 上記した範囲の組成を有する鋼スラブを、 加熱後粗 圧延してシートパーとなし、 該シートパーに仕上圧延を施し、 仕上圧延終了後冷却 して卷き取り熱延板とする熱間圧延工程と、 該熱延板に酸洗および冷間圧延を施し 冷延板とする冷間圧延工程と、 該冷延板に連続焼鈍を行う冷延板焼鈍工程とを、 順 次施すことにより製造される。  The steel sheet of the present invention is basically a steel slab having a composition in the above-mentioned range, which is subjected to rough rolling after heating to form a sheet par. The sheet par is subjected to finish rolling. A hot rolling step of forming a rolled sheet, a cold rolling step of performing pickling and cold rolling on the hot rolled sheet to form a cold rolled sheet, and a cold rolling sheet annealing step of performing continuous annealing on the cold rolled sheet. It is manufactured by sequentially applying.
本発明の製造方法で使用するスラブは、 成分のマクロな偏析を防止すべく連続鎳 造法で製造することが望ましいが、 造塊法、 薄スラブ鏺造法で製造してもよい。 ま た、 スラブを製造した後、 いったん室温まで冷却し、 その後再度加熱する従来法に 加え、 冷却せず温片のままで加熱炉に装入し圧延する直送圧延、 あるいはわずかの 保熱を行った後に直ちに圧延する直接圧延などの省エネルギープロセスも問題なく 適用できる。 とくに、 固溶状態の Nを有効に確保するには直送圧延は有用な技術の 一つである。  The slab used in the production method of the present invention is desirably produced by a continuous production method in order to prevent macroscopic segregation of components, but may be produced by an ingot-making method or a thin slab production method. In addition to the conventional method in which the slab is manufactured and then cooled to room temperature and then heated again, direct-feed rolling, in which the slab is placed in a heating furnace and rolled as it is without cooling, or slight heat retention is performed. Energy saving processes such as direct rolling, in which rolling is performed immediately afterwards, can be applied without any problems. In particular, direct rolling is one of the useful techniques for effectively securing solid solution N.
まず、 熱間圧延工程の条件限定理由について説明する。  First, the reasons for limiting the conditions of the hot rolling step will be described.
スラブ加熱温度： 1000。C以上  Slab heating temperature: 1000. C or more
スラブ加熱温度は、 初期状態として、 必要かつ十分な固溶 N量を確保し、 製品で の固溶 N量の目標値を満足させるために、 1000°C以上とするの好ましい。 なお、 酸 化重量の増加にともなうロスの増大などから 1280。C以下とすることが望ましい。 上記した条件で加熱されたスラブは、 粗圧延によりシートパーとされる。 なお、 粗圧延の条件はとくに規定する必要はなく、 常法にしたがって行えばよい。 しかし、 固溶 N量の確保という観点からはできるだけ短時間で行うのが望ましい。 ついで、 シートバーを仕上げ圧延して熱延板とする。 The slab heating temperature is preferably set to 1000 ° C. or higher in order to secure a necessary and sufficient amount of solute N in the initial state and to satisfy a target value of the amount of solute N in the product. The acid 1280 due to an increase in loss due to an increase in chemical weight. It is desirable to be C or less. The slab heated under the conditions described above is converted into a sheet par by rough rolling. The conditions for rough rolling do not need to be particularly defined, but may be determined according to a conventional method. However, from the viewpoint of securing the amount of dissolved N, it is desirable to carry out the reaction in as short a time as possible. Next, the sheet bar is finish-rolled into a hot-rolled sheet.
なお、 本発明では、 粗圧延と仕上げ圧延の間で、 相前後するシートパー同士を接 合し、 連続的圧延することが望ましい。 接合手段としては、 圧接法でも、 レーザー 溶接法、 電子ビーム溶接法などを用いるのが好ましい。  In the present invention, it is preferable that continuous sheet rolling is performed by joining the adjacent sheet pars between rough rolling and finish rolling. As the joining means, it is preferable to use a laser welding method, an electron beam welding method, or the like even in a pressure welding method.
連続圧延することにより、 コイル （被処理材） の先端および後端のいわゆる圧延 の非定常部がなくなり、 安定した熱延条件がコイル （被処理材） 全長おょぴ全幅に 渡って可能となる。 これは熱延鋼板のみでなく冷延鋼板の断面の形状および寸法を 改善するのに極めて有効である。 また圧延後に、 ホットランテーブル上で冷却する 場合にも常に張力を付与できるため鋼板形状を良好に保つことが可能である。  Continuous rolling eliminates the so-called rolling unsteady portions at the leading and trailing ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length of the coil (material to be processed). . This is extremely effective in improving the cross-sectional shape and dimensions of not only hot-rolled steel sheets but also cold-rolled steel sheets. Further, even when cooling on a hot run table after rolling, tension can always be applied, so that the steel plate shape can be kept good.
また、 連続圧延を行うことでコイル先端を安定して通板できるため、 通常のシー トパーごとの単発圧延では、 通板性おょぴ嚙込み性の問題で適用が難しかった潤滑 圧延を適用することができる。 これにより圧延荷重を低減することができると同時 にロールの面圧をも低減でき、 口ールの寿命延長が可能となる。  In addition, since continuous rolling allows the coil tip to be passed stably, the lubricating rolling, which was difficult to apply due to the problem of threading and penetration, was applied in single-shot rolling for each sheeter. be able to. As a result, the rolling load can be reduced and, at the same time, the surface pressure of the roll can be reduced, and the life of the needle can be extended.
また、 本発明では、 粗圧延と仕上圧延の間の仕上げ圧延機入側で、 シートパーの 幅方向端部を加熱するシートパーエッジヒータ、 シートパーの長さ方向端部を加熱 するシートパーヒータのいずれか一方または両方を使用して、 シートパーの幅方向 およぴ長手方向の温度分布を均一化することが好ましい。 これにより、 鋼板内の材 質ばらつきをさらに小さくすることができる。 シ トパーエッジヒータ、 シートパ 一ヒータは誘導加熱方式のものとするのが好ましい。  Further, according to the present invention, a sheet par edge heater that heats the width direction end of the sheet par and a sheet par heater that heats the length direction end of the sheet par on the side of the finish rolling mill between the rough rolling and the finish rolling are provided. It is preferable to use one or both of them to make the temperature distribution in the width direction and the longitudinal direction of the sheet par uniform. As a result, it is possible to further reduce the material variation in the steel sheet. It is preferable that the sheet edge heater and the sheet heater are of an induction heating type.
使用手順は、 まずシートパーエッジヒータにより幅方向の温度差を補償すること が望ましい。 このときの加熱量は、 鋼組成などにもよるが、 仕上圧延出側での幅方 向温度分布範囲が概ね 20°C以下となるように設定するのが好ましい。 次いでシート パーヒータにより長手方向の温度差を補償する。 このときの加熱量は、 長さ端部温 度が中央部温度よりも概ね 20°C程度高くなるように設定するのが好ましい。 It is desirable to use a sheet-per-edge heater to compensate for the temperature difference in the width direction. The amount of heating at this time depends on the steel composition and the like, but is preferably set so that the temperature distribution range in the width direction at the finish rolling exit side is approximately 20 ° C. or less. Then sheet The temperature difference in the longitudinal direction is compensated by the per heater. The heating amount at this time is preferably set so that the temperature at the end of the length is approximately 20 ° C higher than the temperature at the center.
仕上圧延出側温度： 800 °C以上  Finishing rolling exit side temperature: 800 ° C or more
仕上圧延出側温度 F D Tは、 均一微細な熱延母板組織を得るために、 800 以上 とする。 F D Tが 800 °Cを下回ると、 鋼板の組織が不均一になり、 一部に加工組織 が残留し、 冷延焼鈍工程を経たのちにも、 組織の不均一性が消滅せず残留する。 こ のため、 プレス成形時に種々の不具合を癸生する危険性が増大する。 また、 加工組 織の残留を回避すべく、 高い卷取温度を採用すると、 粗大結晶粒が発生し、 同様の 不具合が発生する。 また、 卷取温度を高温とすることにより、 固溶 N量の顕著な低 下が生ずるため、 目標とする 340MPa以上の引張強さを得ることが困難となる。 この ようなことから、 仕上圧延出側温度 F D Tは 800 以上とした。 さらに機械的性質 を向上させるには、 F D Tを 820 °C以上とすることが望ましい。 なお、 r値の向上 の観点から、 F D Tは Ac 3変態点以上とするのがより好ましい。 またとくに、 F D Tの上限は規定しないが、 過度に高い場合には、 スケール疵などの発生が顕著とな る。 なお、 F D Tは概ね 1000で程度までとするのが好ましい。  The finish rolling exit side temperature FDT is 800 or more in order to obtain a uniformly fine hot-rolled base plate structure. If the FDT is below 800 ° C, the microstructure of the steel sheet becomes non-uniform, the processed microstructure remains in part, and the non-uniform microstructure remains after passing through the cold rolling annealing process. For this reason, there is an increased risk of various defects occurring during press forming. In addition, if a high winding temperature is used to avoid the remaining of the processed tissue, coarse crystals are generated, and the same problem occurs. In addition, when the winding temperature is set to a high temperature, a remarkable decrease in the amount of solute N occurs, so that it is difficult to obtain a target tensile strength of 340 MPa or more. For this reason, the finish rolling exit temperature FDT was set to 800 or more. In order to further improve the mechanical properties, it is desirable that FDT be 820 ° C or more. From the viewpoint of improving the r value, it is more preferable that FDT is equal to or higher than the Ac 3 transformation point. In particular, the upper limit of FDT is not specified, but if it is excessively high, scale flaws and the like will be remarkable. It is preferable that FDT is approximately 1000 or less.
卷取温度： 800 °C以下  Winding temperature: 800 ° C or less
卷取温度 C Tの低下につれて、 鋼板強度が増加する傾向にある。 目標の引張強さ T S 340MPa以上を確保するためには、 C Tは 800 以下とするのが好ましい。 なお、 C Tが 200 で未満では鋼板形状が乱れやすくなり、 実操業上、 不具合を生じる危険 性が高く、 材質の均一性が低下する傾向を示す。 このため、 C Tは 200°C以上とする のが望ましい。 なお、 より材質の均一性が要求される場合には、 C Tは 300 以上 とするのが好ましい。 なお、 より好ましくは 350 以上である。  As the winding temperature CT decreases, the steel sheet strength tends to increase. In order to secure the target tensile strength T S 340 MPa or more, C T is preferably 800 or less. If the CT is less than 200, the shape of the steel sheet tends to be disturbed, and there is a high risk of causing problems in actual operation, and the uniformity of the material tends to decrease. For this reason, CT is desirably 200 ° C or more. When more uniform material is required, CT is preferably 300 or more. The value is more preferably 350 or more.
また、 本発明では、 仕上圧延において、 熱間圧延荷重を低減するために、 潤滑圧 延を行ってもよい。 潤滑圧延を行うことにより、 熱延板の形状 .材質がより均一化 されるという効果がある。 なお、 潤滑圧延の際の摩擦係数は 0. 25〜0. 10の範囲とす るのが好ましい。 また、 潤滑圧延と連続圧延とを組み合わせることによりさらに、 熱間圧延の操業が安定する。 Further, in the present invention, in finish rolling, lubricating rolling may be performed in order to reduce a hot rolling load. By performing lubricating rolling, there is an effect that the shape and material of the hot-rolled sheet are made more uniform. The coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10. In addition, by combining lubrication rolling and continuous rolling, Hot rolling operation is stabilized.
上記した熱間圧延工程を施された熱延板は、 ついで、 冷間圧延工程により、 酸洗 およぴ冷間圧延を施されて冷延板となる。  The hot-rolled sheet that has been subjected to the above-described hot rolling step is then subjected to pickling and cold rolling in a cold-rolling step to become a cold-rolled sheet.
酸洗の条件は通常公知の条件でよく、 とくに限定されない。 なお、 熱延板のスケ ールが極めて薄い場合には、 酸洗を施すことなく直ちに冷間圧延を行ってもよい。 また、 冷間圧延条件は、 通常公知の条件でよく、 とくに限定されない。 なお、 組 織の均一性確保という観点から冷間圧下率は 60%以上とするのが好ましい。 つぎ に、 冷延板焼鈍工程の条件限定理由について説明する。  The conditions for pickling may be generally known conditions, and are not particularly limited. If the scale of the hot rolled sheet is extremely thin, cold rolling may be performed immediately without performing pickling. In addition, the cold rolling conditions may be generally known conditions, and are not particularly limited. In addition, it is preferable that the cold rolling reduction is 60% or more from the viewpoint of ensuring the uniformity of the tissue. Next, the reasons for limiting the conditions of the cold-rolled sheet annealing step will be described.
冷延板は、 ついで連続焼鈍—冷却からなる冷延板焼鈍工程を施される。  The cold rolled sheet is then subjected to a cold rolled sheet annealing step consisting of continuous annealing and cooling.
連続焼鈍温度： フェライ トーオーステナイ トニ相共存域内の温度  Continuous annealing temperature: Temperature in the coexistence region of ferrite toe austenite and toni phase
フェライト一オーステナイ トニ相共存域内の温度で焼鈍することにより、 ァシキ ユラ一フェライ ト相が形成される。 加えて、 フェライト相にも （111 ) 集合組織が 強く発達するため高い r値が得られる。 一方、 フェライ ト一オーステナイ ト二相共 存域を超えてオーステナイ ト単相となる高い温度では、 逆変態と変態により鋼板の 集合組織がランダム化するため r値が低下する。 このため、 本発明では連続焼鈍の 焼鈍温度を再結晶温度以上フェライ トーオーステナイトニ相共存域内の温度に限定 した。 なお、 r値の安定性からオーステナイ トの分率が 10%以上 50%以下となる温 度とするのが好ましい。 また、 連続焼鈍温度が再結晶温度未満では延性が低くなり、 自動車部品用として限定された特殊用途にしか適用できなくなるため、 再結晶温度 以上とすることが好ましい。  Annealing at a temperature in the coexisting region of ferrite-austenite-to-ni phase forms a ferrite-ferrite phase. In addition, a high r value is obtained because the (111) texture develops strongly in the ferrite phase. On the other hand, at high temperatures beyond the ferrite-austenite two-phase coexistence region to become an austenitic single phase, the r-value decreases because the texture of the steel sheet is randomized by reverse transformation and transformation. For this reason, in the present invention, the annealing temperature of the continuous annealing is limited to a temperature not lower than the recrystallization temperature and in the ferrite to austenite dual phase coexisting region. The temperature is preferably set so that the austenite fraction is 10% or more and 50% or less from the viewpoint of the stability of the r value. Further, if the continuous annealing temperature is lower than the recrystallization temperature, the ductility becomes low, so that it can be applied only to special applications limited to automotive parts.
また、 連続焼鈍時間の保持時間は、 生産効率、 組織の微細化、 固溶 N量の確保の 観点から、 できる限り短いほうが好ましい。 操業の安定性の観点から、 保持時間は 1 0 s以上とするのが好ましく、 また、 組織の微細化と固溶 N量の確保という観点から は、 90 s以下とすることが好ましい。 なお、 材質の安定化という観点からは、 20 s 以上とするのがより好ましい。  Further, the holding time of the continuous annealing time is preferably as short as possible from the viewpoints of production efficiency, refining the structure, and securing the amount of solute N. The holding time is preferably at least 10 s from the viewpoint of operation stability, and is preferably at most 90 s from the viewpoint of microstructuring of the structure and securing the amount of dissolved N. In addition, from the viewpoint of stabilizing the material, it is more preferable to set it to 20 s or more.
連続焼鈍後の冷却： 500 °C以下の温度域まで 10〜 300°C/sの冷却速度で冷却 連 続焼鈍における均熱後の冷却は、 組織の微細化、 ァシキユラ一フェライ ト相の形成、 固溶 N量の確保の観点から重要である。 本 明では、 少なくとも、 500 Cooling after continuous annealing: Cooling at a cooling rate of 10 to 300 ° C / s to a temperature range of 500 ° C or less Cooling after soaking in continuation annealing is important from the viewpoint of microstructural refinement, formation of the ferrite-ferrite phase, and securing of the solute N content. In the present invention, at least 500
¾以下の温度域まで lOtVs以上の冷却速度で連続冷却する。 冷却速度が 10°C/s未満 では、 必要量のァシキユラ一フェライ ト量と、 均一でかつ微細な組織と、 十分な量 の固溶 Nを得ることができない。 一方、 冷却速度が 300。C/sを超えると、 鋼板の幅 方向での材質の均一性が不足する。 また、 連続焼鈍後の 10〜 300°C/sの冷却速度で の冷却停止温度が 500 tを超えると、 組織の微細化が達成できない。 調質圧延あ るいはレペラ一加工：伸び率 0. 5〜10% 連 続 Continuously cool at a cooling rate of lOtVs or more to the temperature range below. If the cooling rate is less than 10 ° C / s, it is not possible to obtain the required amount of ferrite, a uniform and fine structure, and a sufficient amount of solute N. On the other hand, the cooling rate is 300. If it exceeds C / s, the uniformity of the material in the width direction of the steel sheet will be insufficient. In addition, if the cooling stop temperature at a cooling rate of 10 to 300 ° C / s after continuous annealing exceeds 500 t, the microstructure cannot be refined. Temper rolling or repeller processing: elongation 0.5 to 10%
本発明では、 冷延焼鈍工程に引き続いて、 形状矯正、 粗度調整の目的で、 調質圧 延またはレべラー加工を施してもよい。 調質圧延あるいはレべラー加工の伸び率が 合計で 0. 5 %未満では、 形状矯正、 粗度調整の所期の目的が達成できない。 一方、 1 0%を超えると、 延性の低下をもたらす。 なお、 5 %以下とすることが延性確保の観 点からより好ましい。 また、 調質圧延とレべラー加工ではその加工形式が相違する が、 その効果は両者で大きな差異がないことを確認している。 調質圧延、 レべラー 加工は、 めっき処理後でも有効である。 第 7の本発明の高張力冷延鋼板の組成限定理由について説明する。  In the present invention, after the cold rolling annealing step, temper rolling or leveling may be performed for the purpose of shape correction and roughness adjustment. If the total elongation of the temper rolling or leveling is less than 0.5%, the intended purposes of shape correction and roughness adjustment cannot be achieved. On the other hand, if it exceeds 10%, the ductility is reduced. Note that the content is more preferably 5% or less from the viewpoint of ensuring ductility. In addition, although the form of temper rolling and leveler processing are different, it has been confirmed that there is no significant difference between the two. Temper rolling and leveling are effective even after plating. The reason for limiting the composition of the high-tensile cold-rolled steel sheet of the seventh invention will be described.
C： 0. 025〜0. 15%  C: 0.025-0.15%
Cは、 鋼板の強度を増加する元素であり、 また、 本発明の重要な構成要件である 組織を均一かつ微細に制御し、マルテンサイ ト相を十分な量確保するため、 0. 025 % 以上含有する必要がある。 一方、 0. 15%を超えると、 鋼板中の炭化物分率が過大と なり、 延性、 さらには成形性が顕著に低下する。 さらにより重要な問題として、 C 含有量が 0. 15%を超えると、 スポット溶接性、 アーク溶接性等が顕著に低下する。 このようなことから、 Cは 0. 025 ~0. 15%の範囲内に限定した。 なお、 成形性の向 上という観点からは、 0. 08%以下とするのが好ましい。 また、 特に良好な延性が要 求される用途では、 0. 05%以下とするのがより好ましい。 Si： 1. 0 %以下 C is an element that increases the strength of the steel sheet, and contains 0.025% or more in order to uniformly and finely control the structure, which is an important component of the present invention, and to secure a sufficient amount of martensite phase. There is a need to. On the other hand, if it exceeds 0.15%, the carbide fraction in the steel sheet becomes excessive, and the ductility and the formability are significantly reduced. More importantly, if the C content exceeds 0.15%, spot weldability, arc weldability, etc. will be significantly reduced. For these reasons, C is limited to the range of 0.025 to 0.15%. From the viewpoint of improving formability, the content is preferably set to 0.08% or less. Further, for applications requiring particularly good ductility, the content is more preferably 0.05% or less. Si: 1.0% or less
Siは、 鋼の延性を顕著に低下させることなく、 鋼板を高強度化させることができ る有用な元素であり、 0. 005 %以上、 より好ましくは 0. 1 %以上含有するのが好ま しい。 一方、 Siは、 熱間圧延時に変態点を大きく上昇させて品質、 形状の確保を困 難にしたり、 あるいはまた表面性状、 化成処理など、 特に鋼板表面の美麗性に悪影 響を与え、 さらにめつき性にも悪影響を及ぼす元素であり、 本発明では 1. 0 %以下 に限定した。 Siが 1. 0 %以下であれば、 上記した悪影響を低く抑えることができる。 なお、強度要求レベルが低く、特に表面の美麗性が要求される用途には、 Siは 0. 5 % 以下とすることが望ましい。  Si is a useful element that can increase the strength of a steel sheet without significantly reducing the ductility of the steel, and is preferably contained at 0.005% or more, more preferably at 0.1% or more. . On the other hand, Si significantly raises the transformation point during hot rolling, making it difficult to ensure quality and shape, or adversely affects the surface properties and chemical treatment, especially the beauty of the steel sheet surface. It is an element that also has an adverse effect on plating, and is limited to 1.0% or less in the present invention. When the content of Si is 1.0% or less, the above-mentioned adverse effects can be suppressed to a low level. In addition, in applications where the strength requirement level is low, and particularly in the case where beautiful surface is required, the content of Si is desirably 0.5% or less.
Mn： 2. 0 %以下  Mn: 2.0% or less
Mnは、 Sによる熱間割れを防止する有効な元素であり、 含有する S量に応じて添 加するのが好ましく、 また Mnは結晶粒の微細化に対し大きな効果があり、 添加して 材質改善に利用することが望ましい。 さらに Mnは、 連続焼鈍後の急速冷却時にマル テンサイトを安定して生成させるために極めて有効な元素である。 Sを安定して固 定するという観点から、 Mnは 0. 2 %以上の含有が望ましい。 また Mnは鋼板強度を増 加させる元素であり、 T S 500MPa超級の強度が要求される場合には 1. 2 %以上含有 するのが望ましい。 なお、 より好ましくは 1. 5 %以上である。  Mn is an effective element for preventing hot cracking due to S. It is preferable to add Mn in accordance with the amount of S contained.Mn has a great effect on refining crystal grains, and Mn is added. It is desirable to use it for improvement. Furthermore, Mn is an extremely effective element for stably generating martensite during rapid cooling after continuous annealing. From the viewpoint of stably fixing S, the content of Mn is desirably 0.2% or more. Mn is an element that increases the strength of the steel sheet, and when a strength of more than 500 MPa in T S is required, it is preferable to contain 1.2% or more. It is more preferably at least 1.5%.
Mn含有量をこのレベルまで高めると、 熱延条件の変動に対する鋼板の機械的性質、 とくに歪時効硬化特性のばらつきが顕著に改善されるという大きな利点がある。 し かし、 Mnを 2. 0 %を超えて過度に含有すると、 本発明の重要な要件の一つである高 r値を得ることが困難となるとともに、 延性が顕著に低下するため、 Mnは 2. 0 %以 下に限定した。 より良好な耐食性と成形性が要求される用途では、 1. 7 %以下とす るのが好ましい。  When the Mn content is increased to this level, there is a great advantage that the mechanical properties of the steel sheet with respect to the fluctuation of the hot rolling conditions, especially the variation of the strain aging hardening characteristics are remarkably improved. However, if Mn is excessively contained in excess of 2.0%, it becomes difficult to obtain a high r value, which is one of the important requirements of the present invention, and ductility is significantly reduced. Is limited to 2.0% or less. For applications requiring better corrosion resistance and moldability, the content is preferably 1.7% or less.
P ： 0. 08%以下  P: 0.08% or less
Pは、 鋼の固溶強化元素として有用な元素であり、 強度増加の観点から 0. 001 %以上、 より好ましくは 0. 015 %以上含有するのが好ましい。 一方、 過度に含有す ると、 鋼を脆化させ、 さらに鋼板の伸びフランジ加工性を悪化させる。 また、 pは 鋼中で偏析する傾向が強いためそれに起因した溶接部の脆化をもたらす。 このため、P is an element useful as a solid solution strengthening element for steel, and preferably contains 0.001% or more, more preferably 0.015% or more, from the viewpoint of increasing strength. On the other hand, excessive This embrittles the steel and further deteriorates the stretch flangeability of the steel sheet. In addition, p has a strong tendency to segregate in steel, which results in embrittlement of the weld. For this reason,
Pは 0. 08%以下に限定した。 なお、 伸ぴフランジ加工性や溶接部靱性がとくに重要 視される用途では Pは 0. 04%以下とするのが好ましい。 P was limited to 0.08% or less. In applications where stretch flangeability and weld toughness are particularly important, P is preferably set to 0.04% or less.
S ： 0. 02%以下  S: 0.02% or less
Sは、 鋼板中では介在物として存在し、 鋼板の延性を減少させ、 さらには耐食性 の劣化をもたらす元素であり、できるだけ低減するのが好ましく、本発明では Sは 0. 02%以下に限定した。 とくに、 良好な加工性が要求される用途には、 Sは 0. 015 % 以下とすることが好ましい。 また、 とくに優れた伸ぴフランジ加工性が要求される 場合には、 Sは 0. 008 %以下とすることが好ましい。 また、 詳細な機構は不明であ るが、 鋼板の歪時効硬化特性を安定して高いレベルに維持するためには、 Sを 0. 008 %以下まで低減するのが有効である。  S is an element that exists as an inclusion in the steel sheet and reduces the ductility of the steel sheet, and furthermore, deteriorates the corrosion resistance.It is preferable that S is reduced as much as possible.In the present invention, S is limited to 0.02% or less. . In particular, for applications requiring good workability, S is preferably set to 0.015% or less. Further, when particularly excellent stretch flangeability is required, S is preferably 0.008% or less. Although the detailed mechanism is unknown, it is effective to reduce S to 0.008% or less in order to stably maintain the strain age hardening property of the steel sheet at a high level.
A1： 0. 02%以下  A1: 0.02% or less
A1は、 脱酸剤として作用し鋼の清浄度を向上させ、 さらには鋼板の組織を微細化 する元素であり、 本発明では 0. 001 %以上の含有が望ましい。 本発明においては、 固溶状態の Nを強化元素として利用するが、 適正範囲の A1を含有したアルミキルド 鋼のほうが、 A1を添加しない従来のリムド鋼に比して、 機械的性質が優れている。 一方、 過剰の A1含有は、 鋼板の表面性状を悪化させ、 さらに固溶状態の Nを顕著に 低下させて、 極めて大きな歪時効硬化量を得ることが困難となる。 このようなこと から、 本発明では A1は 0. 02%以下に限定した。 なお、 材質の安定性という観点から は、 A1は 0, 001 〜0. 015 %とするのが好ましい。 また、 A1含有量の低減は結晶粒の 粗大化につながる懸念もあるが、 本発明では他の合金元素を最適量に制限すること と、 焼鈍条件を最適な範囲とすることにより、 これを有効に防止している。  A1 is an element that acts as a deoxidizing agent, improves the cleanliness of steel, and further refines the structure of the steel sheet. In the present invention, the content of A1 is preferably 0.001% or more. In the present invention, N in a solid solution state is used as a strengthening element.Aluminum-killed steel containing A1 within an appropriate range has better mechanical properties than conventional rimmed steel without A1 added. . On the other hand, excessive A1 content deteriorates the surface properties of the steel sheet, and further significantly reduces N in the solid solution state, making it difficult to obtain an extremely large amount of strain age hardening. For these reasons, A1 is limited to 0.02% or less in the present invention. From the viewpoint of material stability, A1 is preferably set to 0.001 to 0.015%. In addition, there is a concern that the reduction of the A1 content may lead to coarsening of the crystal grains, but in the present invention, it is effective to limit the other alloying elements to the optimum amount and to set the annealing conditions in the optimum range. Has been prevented.
N ： 0. 0050〜0. 0250%  N: 0.0050 to 0.0250%
Nは、 固溶強化と歪時効硬化により鋼板の強度を増加させる元素であり、 本発明 において最も重要な元素である。 また、 本発明では、 適量の Nを含有して、 さらに 上記したように Al含有量を適正値に調整し、 さらには熱延条件、 焼鈍条件などの製 造条件を制御することにより、 冷延製品あるいはめっき製品で必要かつ十分な固溶 状態の Nを確保する。 これにより、 固溶強化と歪時効硬化による強度 （降伏応力お よび引張り強さ） 上昇効果が十分に発揮され、 引張強さ 440MPa以上、 焼付け硬化量 ( B H量） 80MPa 以上、 歪時効処理前後での引張強さの増加量 Δ T S 40MPa以上、 という本発明鋼板の機械的性質の目標値を安定して得ることができる。 N is an element that increases the strength of the steel sheet by solid solution strengthening and strain age hardening, and is the most important element in the present invention. In the present invention, an appropriate amount of N is contained. As described above, by adjusting the Al content to an appropriate value and controlling the manufacturing conditions such as hot rolling conditions and annealing conditions, N in the solid solution state necessary and sufficient for cold rolled products or plated products can be obtained. Secure. As a result, the effect of increasing the strength (yield stress and tensile strength) by solid solution strengthening and strain age hardening is fully exhibited, with a tensile strength of 440 MPa or more, bake hardening amount (BH amount) of 80 MPa or more, and before and after strain aging treatment. , The target value of the mechanical properties of the steel sheet of the present invention, that is, the increase amount of the tensile strength ΔTS of 40 MPa or more can be stably obtained.
Nが 0. 0050%未満では、 上記した強度上昇効果が安定して現れにくい。 一方、 N が 0. 0250%を超えると、 鋼板の內部欠陥癸生率が高くなるとともに、 連続錄造時の スラブ割れなども多発するようになる。 このため、 Nは 0. 0050〜0. 0250%の範囲に 限定した。 なお、 製造工程全体を考慮した材質の安定性 ·歩留り向上という観点か らは、 Nは 0. 0070〜0. 0170%の範囲とするのがより好ましい。 なお、 本発明の範囲 内の N量であれば、 溶接性等への悪影響は全くない。  If N is less than 0.0050%, the above-described effect of increasing the strength is unlikely to appear stably. On the other hand, if N exceeds 0.0250%, the rate of partial cracking of the steel sheet will increase, and slab cracking will occur more frequently during continuous forming. For this reason, N was limited to the range of 0.0050 to 0.0250%. It is more preferable that N is in the range of 0.0070% to 0.0170% from the viewpoints of material stability and yield improvement in consideration of the entire manufacturing process. If the N content is within the range of the present invention, there is no adverse effect on weldability and the like.
固溶状態の N： 0. 0010%以上  N in solid solution state: 0.0010% or more
冷延製品で十分な強度が確保され、 さらに Nによる歪時効硬化が有効に発揮され るには、 鋼板中に固溶状態の N (固溶 Nともいう） が少なく とも 0. 0010%以上存在 する必要がある。  In order for a cold-rolled product to have sufficient strength and for strain age hardening to be effectively exhibited by N, at least 0.0010% or more of N in solid solution (also called solid solution N) is present in the steel sheet There is a need to.
ここで、 固溶 N量は、 鋼中の全 N量から、 析出 N量を差し引いた値を固溶 Nとす る。 なお、 析出 N量の分析法としては、 本発明者らが種々の方法を比較検討した結 果、 定電位電解法を用いた電解抽出分析法により求めるのが有効である。 なお、 抽 出分析に用いる地鉄を溶解する方法として、 酸分解法、 ハロゲン法おょぴ電解法が ある。 この中で、 電解法は炭化物、 窒化物などの極めて不安定な析出物を分解する ことなく、 安定して地鉄のみを溶解できる。 電解液としては、 ァセチル . アセトン 系を用いて、 定電位にて電解する。 本発明では定電位電解法を用いて析出 N量を測 定した結果が、 実際の材質の変化とよい対応を示した。  Here, the solid solution N amount is defined as a value obtained by subtracting the precipitated N amount from the total N amount in the steel. As a method of analyzing the amount of precipitated N, it is effective to obtain the amount by the electrolytic extraction analysis method using the potentiostatic electrolysis method as a result of comparative studies of various methods by the present inventors. In addition, there are acid decomposition method, halogen method and electrolysis method as a method for dissolving base iron used for extraction analysis. Among them, the electrolysis method can stably dissolve only ground iron without decomposing extremely unstable precipitates such as carbides and nitrides. Electrolyte at a constant potential using an acetyl.acetone system as the electrolytic solution. In the present invention, the result of measuring the amount of deposited N using the potentiostatic electrolysis method showed a good correspondence with the actual material change.
このようなことから、 本発明では、 定電位電解法により抽出した残渣を化学分析 して残渣中の N量を求め、 これを析出 N量とする。 なお、 さらに高い BH量、 ATSが必要な場合には、 固溶 N量を 0.0020%以上、 さらに高い値を得るためには、 0.0030%以上とするのが好ましい。 固溶 N量の上限 値は特に限定しないが、 添加した全 N量がすべて残留しても機械的性質の低下は小 さい。 For this reason, in the present invention, the residue extracted by the potentiostatic electrolysis is subjected to chemical analysis to determine the amount of N in the residue, which is defined as the amount of deposited N. When a higher BH content and ATS are required, the amount of solute N is preferably 0.0020% or more. To obtain a higher value, the content is preferably 0.0030% or more. The upper limit of the amount of solute N is not particularly limited, but even if all of the added N remains, the decrease in mechanical properties is small.
NZA1 (N含有量と A1含有量の比） ： 0.3 以上  NZA1 (Ratio of N content and A1 content): 0.3 or more
製品状態で、 固溶 Nを 0.0010%以上安定させて残留させるためには、 Nを強力に 固定する元素である A1の量を制限する必要がある。 本発明の組成範囲内の N含有量 (0.0050〜0.0250%) と A1含有量 （0.02%以下） の組合せを広範囲に変えた鋼板に ついて検討した結果、 NZA1を 0.3 以上とすることにより、 冷延製品おょぴめっき 製品での固溶 Nを安定して 0.0010%以上とすることができることがわかった。 この ため、 NZA1を 0.3 以上に限定した。  In the product state, it is necessary to limit the amount of A1, which is an element that strongly fixes N, in order to stably retain solute N of 0.0010% or more in the product state. As a result of examining a steel sheet in which the combination of the N content (0.0050 to 0.0250%) and the A1 content (0.02% or less) in the composition range of the present invention was changed over a wide range, the cold rolling was performed by setting NZA1 to 0.3 or more. Product plating It was found that the solid solution N in the product can be stably increased to 0.0010% or more. Therefore, NZA1 was limited to 0.3 or more.
本発明では、 上記した組成に加えてさらに、 次 d群〜 g群  In the present invention, in addition to the above composition, the following d group to g group
d群： Cu、 Ni、 Cr、 Moの 1種または 2種以上を合計で 1.0 %以下  Group d: 1% or more of Cu, Ni, Cr, Mo, 1.0% or less in total
e群： Nb、 Ti、 Vの 1種または 2種以上を合計で 0.1 %以下  Group e: 0.1% or less in total of one or more of Nb, Ti, and V
f 群： Bを 0.0030%以下  Group f: 0.0030% or less for B
g群： Ca、 REM の1種または2種を合計で0.0010〜0.010. %  g group: Ca or REM 1 or 2 kinds in total 0.0010 to 0.010.%
の 1群または 2群以上を含有するのが好ましい。 It is preferable that one or two or more groups are contained.
d群の元素： Cu、 Ni、 Cr、 Moは、 いずれも鋼板の強度上昇に寄与する元素であり、 必要に応じ選択して単独または複合して含有できる。 このような効果は、 Cu、 Ni、 C r、 Moをそれぞれ 0.005 %以上の含有で認められる。 しかし、 含 量が多すぎると熱 間変形抵抗が増加し、 あるいは化成処理性や広義の表面処理特性が悪化するうえ、 溶接部が硬化し溶接部成形性が劣化する。 また r値も低下する傾向がある。 このた め、 a群の元素は合計で 1.0 %以下とするのが好ましい。 なお、 Moは、 0.05%以上 多量に含有すると顕著に r値を低下させる場合があり、 本発明では Moを含有する場 合は 0.05%未満に限定するのが好ましい。  Group d elements: Cu, Ni, Cr, and Mo are all elements that contribute to the increase in the strength of the steel sheet, and can be selected as necessary and contained alone or in combination. Such an effect is recognized when the content of Cu, Ni, Cr, and Mo is 0.005% or more, respectively. However, when the content is too large, the hot deformation resistance increases, or the chemical conversion property and the surface treatment properties in a broad sense deteriorate, and the welded part is hardened and the weldability is deteriorated. Also, the r value tends to decrease. For this reason, it is preferable that the total of the elements in group a be 1.0% or less. When Mo is contained in a large amount of 0.05% or more, the r value may be significantly reduced. In the present invention, when Mo is contained, the content is preferably limited to less than 0.05%.
e群の元素： Nb、 Ti、 Vは、 いずれも結晶粒の微細化 ·均一化に寄与する元素で あり、 必要に応じ選択して単独または複合して含有できる。 このような効果は、 Nb、 Ti、 Vをそれぞれ 0. 005 %以上の含有で認められる。 しかし、 含有量が多すぎると、 熱間変形抵抗が増加し、 また化成処理性や広義の表面処理特性が悪化する。 このた め、 b群の元素は合計で 0. 1 %以下とするのが好ましい。 Elements of group e: Nb, Ti, and V are elements that contribute to the refinement and uniformity of crystal grains. Yes, and can be selected alone or in combination as needed. Such an effect is observed when Nb, Ti, and V are each contained at 0.005% or more. However, if the content is too large, the hot deformation resistance increases, and the chemical conversion property and the surface treatment properties in a broad sense deteriorate. For this reason, it is preferable that the total of the elements of group b be 0.1% or less.
f 群の元素： Bは、 鋼の焼入れ性を向上させる効果を有する元素であり、 フェラ ィ ト相以外の低温変態相の分率を増加させて、 鋼の強度を増加させる目的で必要に 応じ含有することができる。 このような効果は、 Bを 0. 0005%以上の含有で認めら れる。 しかし、 量が多すぎると熱間変形能が低下し、 BNを生成することで固溶 Nを 低減させる。 このため、 Bは 0. 0030%以下とするが好ましい。  Elements of group f: B is an element that has the effect of improving the hardenability of steel, and increases the fraction of low-temperature transformation phases other than the ferrite phase, as necessary to increase the strength of steel. Can be contained. Such an effect is recognized when B is contained at 0.0005% or more. However, if the amount is too large, the hot deformability decreases, and BN is formed to reduce the solute N. Therefore, B is preferably set to 0.0030% or less.
g群の元素： Ca、 REM は、 いずれも介在物の形態制御に役立つ元素であり、 特に 伸ぴフランジ成形性の要求がある場合には、 単独または複合して含有するのが好ま しい。 その場合、 d群の元素の合計で、 0. 0010%未満では介在物の形態制御効果が 不足し、 一方、 0. 010 %を超えると表面欠陥の発生が目立つようになる。 このため、 d群の元素を合計で 0. 0010~0. 010 %の範囲に限定することが好ましく、 これによ り、 表面欠陥の発生を伴うことなく伸ぴフランジ加工性を改善することができる。 つぎに、 本発明鋼板の組織について説明する。  Elements of group g: Ca and REM are both elements that are useful for controlling the morphology of inclusions, and particularly when stretch flangeability is required, it is preferable to include them alone or in combination. In this case, if the total of the elements in group d is less than 0.0010%, the effect of controlling the morphology of inclusions is insufficient, while if it exceeds 0.010%, the occurrence of surface defects becomes noticeable. For this reason, it is preferable to limit the total number of elements in the d group to the range of 0.0010 to 0.010%, thereby improving the stretch flange workability without generating surface defects. it can. Next, the structure of the steel sheet of the present invention will be described.
フェライ ト相の面積率： 80%以上  Area ratio of ferrite phase: 80% or more
本発明の冷延鋼板は、 ある程度の加工性が要求される自動車用鋼板を対象として おり、 延性を確保するために、 フ ライ ト相を面積率で 80%以上含む組織とする。 フェライ ト相の面積率が 80%未満では、 加工性が要求される自動車用鋼板として必 要な延性を確保することが困難となる。 なお、 さらに良好な延性が要求される場合 は、 フェライ ト相の面積率は 85%以上とするのが望ましい。 なお、 本発明でいうフ ェライ トは、 歪みが残留していない状態のいわゆるポリゴナルなフェライ トをいう ものとする。  The cold-rolled steel sheet according to the present invention is intended for a steel sheet for automobiles that requires a certain degree of workability, and has a structure containing 80% or more of a fly phase in area ratio in order to ensure ductility. If the area ratio of the ferrite phase is less than 80%, it will be difficult to secure the required ductility as an automotive steel sheet that requires workability. If even better ductility is required, the area ratio of the ferrite phase should be 85% or more. The ferrite according to the present invention refers to a so-called polygonal ferrite in which no distortion remains.
フェライ ト相の平均結晶粒径： 10 /x m 以下  Average grain size of ferrite phase: 10 / x m or less
本発明では平均結晶粒径として、 断面組織写真から A S TMに規定された求積法 により算出した値と、同じく A S TMに規定された切断法により求めた公称粒径（例 えば梅本ら ：熱処理， 24 (1984) ，334参照） のうち、 より大きい方を採用する。 本発明の冷延鋼板では、 製品段階で所定量の固溶 N量を確保しているが、 本発明 者らの実験 ·検討によれば、 同一量の固溶 Nを有する鋼板でも、 歪時効硬化特性に ばちつきが生じる場合があり、 その主たる要因の一つが結晶粒径であることが判明 した。 平均結晶粒径を少なくとも 10 /z ra 以下、 望ましくは 8 /χ πι 以下にすることに より、 安定して高い B H量、 A T Sが得られる。 詳細な機構は不明であるが、 結晶 粒界への合金元素の偏祈と析出、 さらにはこれらに及ぼす加工、 熱履歴の影響に関 係していると推定される。 In the present invention, the average crystal grain size is calculated by the quadrature The larger of the value calculated by the above and the nominal particle size (see, for example, Umemoto et al .: Heat treatment, 24 (1984), 334) determined by the cutting method also specified in ASTM is adopted. In the cold-rolled steel sheet of the present invention, a predetermined amount of solute N is ensured at the product stage. However, according to experiments and studies conducted by the present inventors, even in a steel sheet having the same amount of solute N, strain aging has occurred. In some cases, the hardening characteristics may flicker, and it has been found that one of the main factors is the crystal grain size. By setting the average grain size to at least 10 / z ra or less, preferably 8 / χπι or less, a stable high BH content and ATS can be obtained. Although the detailed mechanism is unknown, it is presumed to be related to the bias and precipitation of alloying elements at the grain boundaries, and the effects of processing and thermal history on these.
したがって、 歪時効硬化特性の安定化を図るためには、 フェライ ト相の平均結晶 粒径を ΙΟ μ ηι 以下、 好ましくは 8 /ζ ιη 以下とする必要がある。  Therefore, in order to stabilize the strain age hardening characteristics, the average crystal grain size of the ferrite phase needs to be と す る μηι or less, preferably 8 / ζιη or less.
以上のように自動車用鋼板としての延性を確保し、 かつ歪時効硬化特性の安定化 を図るため、 本発明では平均結晶粒径 10 / Di 以下のフェライ トを面積率で 80%以上 含む組織とする。  As described above, in order to ensure ductility as an automotive steel sheet and stabilize the strain aging hardening characteristics, the present invention provides a structure containing 80% or more of ferrite having an average crystal grain size of 10 / Di or less by area ratio. I do.
マルテンサイ ト相の面積率： 2 %以上  Martensite phase area ratio: 2% or more
本発明の冷延鋼板は、 第 2相として、 マルテンサイ ト相を面積率で 2 %以上含有 する。 マルテンサイ ト相が 2 %以上存在することにより、 良好な延性と、 さらに、 大きな歪時効硬化量が得られる。 詳細な機構は不明であるが、 マルテンサイ ト相の 存在により、 時効前の予歪み加工時に極めて有効に歪が内部に蓄積されるためと推 定される。 さらに、 マルテンサイ ト相の存在は、 時効劣化を改善するためにも有効 である。 なお、 良好な強度一延性パランス、 低降伏比を得るためには、 マルテンサ イト相の面積率を 5 °/0以上とするのが好ましい。 なお、 20%を超える多量のマルテ ンサイ ト相の存在は延性の低下という問題がある。 このため、 マルテンサイ ト相の 面積率は、 2 %以上、 好ましくは 5 %以上、 20%以下である。 The cold-rolled steel sheet of the present invention contains, as the second phase, a martensite phase in an area ratio of 2% or more. When the martensite phase is present in an amount of 2% or more, good ductility and a large amount of strain age hardening can be obtained. Although the detailed mechanism is unknown, it is presumed that the presence of the martensite phase causes the strain to accumulate extremely effectively during prestraining before aging. Furthermore, the presence of the martensite phase is also effective in improving aging degradation. In order to obtain a good strength-ductility balance and a low yield ratio, the area ratio of the martensite phase is preferably 5 ° / 0 or more. The presence of a large amount of martensite phase exceeding 20% has a problem that ductility is reduced. Therefore, the area ratio of the martensite phase is 2% or more, preferably 5% or more and 20% or less.
第 2相として、 上記したマルテンサイ ト相以外に、 パーライト、 べィナイ ト、 残 留オーステナイトが存在することはなんら問題はないが、 本発明ではフェライ ト相 分率を 80%以上マルテンサイ ト相分率を 2 %以上とする必要があり、 パーライ ト、 ペイナイ ト、 残留オーステナイ トの合計の面積率で 18%未満に限定される。 As the second phase, there is no problem that pearlite, bainite and residual austenite exist in addition to the above-mentioned martensite phase, but in the present invention, the ferrite phase The fraction must be 80% or more and the martensite phase fraction must be 2% or more. The total area ratio of perlite, payite and residual austenite is limited to less than 18%.
上記した組成と組織を有する本発明の冷延鋼板は、 引張強さ （T S ) 440 MPa 以上で概ね 780MPa以下を有し、さらに母相フェライトの集合組織制御により r値が 1. 2 以上の高 r値と、 優れた歪時効硬化特性を有する冷延鋼板である。 T Sが 440MPa を下回る鋼板では、 構造部材的な要素をもつ部材に広く適用することができない。 また、 さらに適用範囲を拡げるには T Sは 500MPa以上とするのが望ましい。 また、 r値が 1. 2 未満では、 広範囲なプレス成形部品に適用できない。 なお、 r値の好ま しい範囲は 1. 4以上である。  The cold-rolled steel sheet of the present invention having the above-described composition and structure has a tensile strength (TS) of 440 MPa or more and about 780 MPa or less, and further has a high r-value of 1.2 or more by controlling the texture of the matrix ferrite. It is a cold-rolled steel sheet that has an r-value and excellent strain aging hardening characteristics. Steel sheets with TS below 440MPa cannot be widely applied to members with structural elements. To further expand the application range, it is desirable that T S be 500 MPa or more. If the r value is less than 1.2, it cannot be applied to a wide range of press-formed parts. The preferred range of the r value is 1.4 or more.
本発明において 「優れた歪時効硬化特性」 とは、 上記したように、 引張歪 5 %の 予変形後、 170 の温度に 20min 保持する条件で時効処理したとき、 この時効処理 前後の変形応力増加量 （B H量と記す； B H量 =時効処理後の降伏応力一時効処理 前の予変形応力） が 80MPa 以上であり、 かつ歪時効処理 （前記予変形 +前記時効処 理） 前後の引張強さ増加量 （A T Sと記す； A T S =時効処理後の引張強さ一予変 形前の引張強さ） が 40MPa 以上であることを意味する。  In the present invention, “excellent strain aging hardening characteristics” means that, as described above, after pre-deformation with a tensile strain of 5%, when subjected to aging treatment at a temperature of 170 for 20 minutes, the deformation stress before and after this aging treatment is increased. Amount (denoted as BH amount; BH amount = pre-deformation stress before yield stress temporary treatment after aging treatment) is 80MPa or more and tensile strength before and after strain aging treatment (pre-deformation + aging treatment) It means that the amount of increase (ATS; ATS = tensile strength after aging treatment-tensile strength before pre-deformation) is 40 MPa or more.
歪時効硬化特性を規定する場合、 予歪 （予変形） 量が重要な因子となる。 本発明 者らは、 自動車用鋼板に適用される変形様式を想定して、 歪時効硬化特性に及ぼす 予歪量の影響について調査し、 その結果、 ①前記変形様式における変形応力は、 極 めて深い絞り加工の場合を除き、 概ね 1軸相当歪 （引張歪） 量で整理できること、 ②実部品ではこの 1軸相当歪量が概ね 5 %を上回っていること、 ③部品強度が、 予 歪 5 %の歪時効処理後に得られる強度 （Y Sおよび T S ) と良く対応することを突 き止めた。 この知見をもとに、 本発明では、 歪時効処理の予変形を引張歪 5 %に定 めた。  When defining the strain age hardening characteristics, the amount of prestrain (prestrain) is an important factor. The present inventors have investigated the effect of the amount of pre-strain on the strain aging hardening characteristics, assuming a deformation mode applied to a steel sheet for automobiles. As a result, (1) the deformation stress in the above-mentioned deformation mode is extremely low. Except in the case of deep drawing, it can be roughly organized by the strain (tensile strain) equivalent to one axis. (2) In actual parts, the strain equivalent to one axis is more than about 5%. It has been found that the strength (YS and TS) obtained after the strain aging treatment of% corresponds well. Based on this finding, in the present invention, the pre-strain of the strain aging treatment was determined to be 5% tensile strain.
従来の塗装焼付け処理条件は、 170 °C X 20min が標準として採用されている。 な お、 多量の固溶 Nを含む本発明鋼板に 5 %以上の歪が加わる場合は、 より緩やかな (低温側の） 処理でも硬化が達成され、 言い換えれば時効条件をより幅広くとるこ とが可能である。 また、 一般に、 硬化量を稼ぐには、 過度の時効で軟化させない限 りにおいて、 より高温で、 より長時間保持することが有利である。 The conventional paint baking condition is 170 ° C X 20 min as standard. If a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of solute N, hardening is achieved even with milder (lower temperature) treatment, in other words, aging conditions can be broadened. And it is possible. In general, in order to increase the amount of hardening, it is advantageous to hold at a higher temperature and for a longer time, unless softening is caused by excessive aging.
具体的に述べると、 本発明鋼板では、 予変形後に硬化が顕著となる加熱温度の下 限は概ね 100 である。 一方、 加熱温度が 300 を超えると硬化が頭打ちとなり、 逆にやや軟化する傾向が現れるほか、 熱歪やテンパーカラーの発生が目立つように なる。 また、 保持時間については、 加熱温度 200 で程度のとき概ね 30 s程度以上と すれば略十分な硬化が達成される。 さらに大きな安定した硬化を得るには保持時間 6 0 s以上とするのが好ましい。 しかし、 20min を超える保持では、 さらなる硬化を望 みえないばかり力 生産効率も著しく低下して実用面では不利である。  Specifically, in the steel sheet of the present invention, the lower limit of the heating temperature at which hardening is remarkable after pre-deformation is approximately 100. On the other hand, when the heating temperature exceeds 300, curing hardens, and on the contrary, it tends to soften slightly, and the occurrence of heat distortion and temper color becomes conspicuous. If the holding time is about 30 s or more when the heating temperature is about 200, almost sufficient curing can be achieved. In order to obtain even greater stable curing, the holding time is preferably 60 s or more. However, holding for more than 20 min is not practical because it does not allow further hardening and significantly reduces power production efficiency.
以上のことから、 本発明では、 時効処理条件として従来の塗装焼付処理条件の加 熱温度である 170 で、 保持時間を 20min で評価すると定めた。 従来の塗装焼付け型 鋼板では十分な硬化が達成されない低温加熱 ·短時間保持の時効処理条件下でも、 本発明鋼板では大きな硬化が安定的に達成される。 なお、 加熱の仕方はとくに制限 されず、 通常の塗装焼付けに採用されている炉による雰囲気加熱のほか、 たとえば 誘導加熱や、 無酸化炎、 レーザ、 プラズマなどによる加熱などのいずれも好ましく 用いうる。  From the above, in the present invention, it was determined that the holding time was evaluated at 170 min, which is the heating temperature under the conventional paint baking treatment conditions, and the holding time was evaluated at 20 min as the aging treatment conditions. Even under the aging condition of low-temperature heating and short-time holding, in which sufficient hardening is not achieved with the conventional paint-baked steel sheet, large hardening is stably achieved in the steel sheet of the present invention. The method of heating is not particularly limited, and in addition to atmospheric heating using a furnace employed for ordinary coating baking, for example, induction heating, heating using a non-oxidizing flame, laser, plasma, or the like can be preferably used.
自動車用の部品強度は外部からの複雑な応力負荷に抗しうる必要があり、 それゆ え素材鋼板では小さな歪域での強度特性だけでなく大きな歪域での強度特性も重要 となる。 本発明者らはこの点に鑑み、 自動車部品の素材となすべき本発明鋼板の B H量を 80MPa 以上とするとともに、 Δ T S量を 40MPa 以上とする。 なお、 より好ま しくは、 B H量 lOOMPa以上、 Δ T S 50MPa以上とする。 また、 時効処理の際の加熱 温度をより高温側に、 および または、 保持時間をより長時間側に、 設定すること により、 B H量、 A T S量をより大きくすることができる。  The strength of automotive components needs to be able to withstand complex external stress loads, so that not only the strength characteristics in a small strain range but also the strength characteristics in a large strain range are important for a material steel plate. In view of this point, the present inventors set the BH amount of the steel sheet of the present invention, which is to be used as a material for automobile parts, to be 80 MPa or more and the ΔTS amount to be 40 MPa or more. More preferably, the amount of BH should be lOOMPa or more, and ΔTS should be 50MPa or more. In addition, by setting the heating temperature during the aging treatment to a higher temperature side and / or the holding time to a longer time side, the BH amount and the ATS amount can be further increased.
また、 本発明鋼板は、 成形後、 とくに加熱を行なわず、 室温で 1週間程度放置し ておくだけで、 完全時効の 40%程度の強度の增加が期待できるという利点がある。 また、 本発明鋼板は、 成形加工されない状態では、 室温で長時間放置されても時 効劣化 （Y Sが増加しかつ E l (伸び） が減少する現象） は起こらないという、 従 来の時効性鋼板にない利点も備わっている。 なお、 実際のプレス成形で不具合を生 じないためには、 プレス成形前の室温における 3か月間の時効で、 Y Sの増加量が 3 OMPa 以下、 伸びの低下が 2 %以下、 降伏点伸びの回復が 0. 2 %以下となることが必 要となる。 In addition, the steel sheet of the present invention has an advantage that it can be expected to increase the strength by about 40% of the full aging simply by leaving it at room temperature for about one week without heating after forming. In addition, the steel sheet of the present invention, even if left untreated at room temperature for a long time, It also has the advantage that conventional aging steel sheets do not have deterioration effect (phenomenon of increasing YS and decreasing El (elongation)), which is not possible with conventional aging steel sheets. In order to prevent failures in the actual press forming, aging at room temperature for 3 months before press forming increases the YS by 3 OMPa or less, decreases the elongation by 2% or less, and increases the yield point elongation. Recovery must be less than 0.2%.
また、 本発明では、 上記した本発明冷延鋼板の表面に電気めつきまたは溶融めつ きを施しても何ら問題はない。 これらめつき鋼板も、 めっき前と同程度の T S、 B H量、 A T S量を示す。 めっきの種類としては、 電気亜鉛めつき、 溶融亜鉛めつき、 合金化溶融亜鉛めつき、 電気錫めつき、 電気クロムめつき、 電気ニッケルめっき等、 いずれも好ましく適用しうる。 第 8の本発明による鋼板の製造方法について説明する。  Further, in the present invention, there is no problem even if the surface of the above-mentioned cold rolled steel sheet of the present invention is electroplated or melted. These plated steel sheets also show the same amount of T S, B H, and A T S as before plating. As the type of plating, any of electroplating, hot-dip galvanizing, alloyed hot-dip galvanizing, electro-tin plating, electro-chrome plating, and electro-nickel plating can be preferably applied. An eighth method of manufacturing a steel sheet according to the present invention will be described.
本発明鋼板は、 基本的には、 上記した範囲の組成を有する鋼スラブを、 加熱後粗 圧延してシートパーとなし、 該シートパーに仕上圧延を施し、 仕上圧延終了後冷却 して卷き取り熱延板とする熱間圧延工程と、 該熱延板に酸洗およぴ冷間圧延を施し 冷延板とする冷間圧延工程と、 該冷延板に箱焼鈍を施し、 ついで連続焼鈍を行う冷 延板焼鈍工程とを、 順次施すことにより製造される。  The steel sheet of the present invention is basically a steel slab having a composition in the above-mentioned range, which is subjected to rough rolling after heating to form a sheet par. The sheet par is subjected to finish rolling. A hot rolling step of forming a rolled sheet, a pickling and cold rolling of the hot rolled sheet to form a cold rolled sheet, a box annealing of the cold rolled sheet, and a continuous annealing It is manufactured by sequentially performing a cold rolled sheet annealing step to be performed.
本発明の製造方法で使用するスラブは、 成分のマクロな偏析を防止すべく連続鎳 造法で製造することが望ましいが、 造塊法、 薄スラブ鎳造法で製造してもよい。 ま た、 スラブを製造した後、 いったん室温まで冷却し、 その後再度加熱する従来法に 加え、 冷却せず温片のままで加熱炉に装入し圧延する直送圧延、 あるいはわずかの 保熱を行った後に直ちに圧延する直接圧延などの省エネルギープロセスも問題なく 適用できる。 とくに、 固溶状態の Nを有効に確保するには直送圧延は有用な技術の 一^ 3である。  The slab used in the production method of the present invention is desirably produced by a continuous production method in order to prevent macroscopic segregation of components, but may be produced by an ingot-making method or a thin slab production method. In addition to the conventional method in which the slab is manufactured and then cooled to room temperature and then heated again, direct-feed rolling, in which the slab is placed in a heating furnace and rolled as it is without cooling, or slight heat retention is performed. Energy saving processes such as direct rolling, in which rolling is performed immediately afterwards, can be applied without any problems. In particular, direct rolling is one of the most useful technologies for effectively securing solid solution N.
まず、 熱間圧延工程の条件限定理由について説明する。  First, the reasons for limiting the conditions of the hot rolling step will be described.
スラブ加熱温度： 1000°C以上 スラブ加熱温度は、 熱間圧延の初期状態として、 必要かつ十分な固溶 N量を確保 し、 製品での固溶 N量は目標値を満足させるために、 1000°C以上とするの好ましい。 なお、 酸化重量の増加にともなうロスの増大などから 1280°C以下とすることが望ま しい。 Slab heating temperature: 1000 ° C or more The slab heating temperature is preferably at least 1000 ° C in order to secure a necessary and sufficient amount of solute N in the initial state of hot rolling and to satisfy the target value of solute N in the product. It is desirable to set the temperature to 1280 ° C or less because of the increase in loss due to the increase in oxidation weight.
上記した条件で加熱されたスラブは、 粗圧延によりシートパーとされる。 なお、 粗圧延の条件はとくに規定する必要はなく、 常法にしたがって行えばよい。 しかし、 固溶 N量の確保という観点からはできるだけ短時間で行うのが望ましい。 ついで、 シートパーを仕上げ圧延して熱延板とする。  The slab heated under the conditions described above is converted into a sheet par by rough rolling. The conditions for rough rolling do not need to be particularly defined, but may be determined according to a conventional method. However, from the viewpoint of securing the amount of dissolved N, it is desirable to carry out the reaction in as short a time as possible. Next, the sheet par is finish-rolled into a hot-rolled sheet.
なお、 本発明では、 粗圧延と仕上げ圧延の間で、 相前後するシートパー同士を接 合し、 連続圧延することが望ましい。 接合手段としては、 圧接法でも、 レーザー溶 接法、 電子ビーム溶接法などを用いるのが好ましい。  In the present invention, it is preferable that continuous sheet rolling is performed by joining the adjacent sheet pars between the rough rolling and the finish rolling. As a joining means, it is preferable to use a laser welding method, an electron beam welding method, or the like even in a pressure welding method.
連続圧延することにより、 コイル （被処理材） の先端おょぴ後端のいわゆる圧延 の非定常部がなくなり、 安定した熱延条件がコイル （被処理材） 全長および全幅に 渡って可能となる。 これは熱延鋼板のみでなく冷延鋼板の断面の形状および寸法を 改善するのに極めて有効である。 また圧延後に、 ホットランテーブル上で冷却する 場合にも常に張力を付与できるため鋼板形状を良好に保つことが可能である。  Continuous rolling eliminates the so-called unsteady rolling part at the front and rear ends of the coil (material to be processed), and enables stable hot rolling conditions over the entire length and width of the coil (material to be processed). . This is extremely effective in improving the cross-sectional shape and dimensions of not only hot-rolled steel sheets but also cold-rolled steel sheets. Further, even when cooling on a hot run table after rolling, tension can always be applied, so that the steel plate shape can be kept good.
また、 連続圧延を行うことでコイル先端を安定して通板できるため、 通常のシー トバーごとの単発圧延では、 通板性およぴ嚙込み性の問題で適用できなかった潤滑 圧延を適用することができる。 これにより圧延荷重を低減することができると同時 にロールの面圧をも低減でき、 ロールの寿命延長が可能となる。  In addition, since continuous rolling enables stable coil passing at the coil end, lubricating rolling, which could not be applied due to problems with threading and penetration, was applied in single-shot rolling for each sheet bar. be able to. As a result, the rolling load can be reduced and, at the same time, the surface pressure of the roll can be reduced, and the life of the roll can be extended.
また、 本発明では、 粗圧延と仕上圧延の間の仕上げ圧延機入側で、 シートパーの 幅端部を加熱するシートパーエッジヒータ、 シートパーの長さ端部を加熱するシー トパーヒータのいずれか一方または両方を使用して、 シートパーの幅方向おょぴ長 手方向の温度分布を均一化することが好ましい。 これにより、 鋼板内の材質ばらつ きをさらに小さくすることができる。 シートパーエッジヒータ、 シートバーヒータ は誘導加熱方式のものとするのが好ましい。 使用手順は、 まずシートパーエッジヒータにより幅方向の温度差を補償すること が望ましい。 このときの加熱量は、 鋼組成などにもよるが、 仕上圧延出側での幅方 向温度分布範囲が概ね 20°C以下となるように設定するのが好ましい。 次いでシート パーヒータにより長手方向の温度差を補償する。 このと.きの加熱量は、 長さ端部温 度が中央部温度よりも概ね 20°C程度高くなるように設定するのが好ましい。 Further, according to the present invention, at the entry side of the finishing mill between the rough rolling and the finish rolling, one or both of a sheet par edge heater for heating the width end of the sheet par and a sheet par heater for heating the length end of the sheet par are provided. It is preferable to use both of them to equalize the temperature distribution in the width direction of the sheet par. Thereby, the variation in the material in the steel sheet can be further reduced. The sheet-per-edge heater and the sheet-bar heater are preferably of an induction heating type. It is desirable to use a sheet-per-edge heater to compensate for the temperature difference in the width direction. The amount of heating at this time depends on the steel composition and the like, but is preferably set so that the temperature distribution range in the width direction at the finish rolling exit side is approximately 20 ° C. or less. Next, the temperature difference in the longitudinal direction is compensated for by the sheet heater. It is preferable that the heating amount in this case is set so that the temperature at the end of the length is approximately 20 ° C higher than the temperature at the center.
仕上圧延出側温度： 800 以上  Finishing rolling exit temperature: 800 or more
仕上圧延出側温度 F D Tは、 均一微細な熱延母板組織を得るために、 800 °C以上 とする。 F D Tが 800 °Cを下回ると、 鋼板の組織が不均一になり、 一部に加工組織 が残留し、 冷延焼鈍工程を経たのちにも、 組織の不均一性が消滅せず残留する。 こ のため、 プレス成形時に種々の不具合を発生する危険性が增大する。 また、 加工組 織の残留を回避すべく、 高い卷取温度を採用すると、 粗大結晶粒が発生し、 同様の 不具合が発生する。 また、 卷取温度を高温とすることにより、 固溶 N量の顕著な低 下が生ずるため、 目標とする 440MPa以上の引張強さを得ることが困難となる。 この ようなことから、 仕上圧延出側温度 F D Tは 800 °C以上とした。 さらに機械的性質 を向上させるには、 F D Tを 820 °C以上とすることが望ましい。 とくに、 F D Tの 上限は規定しないが、 過度に高い場合には、 スケール疵などの発生が顕著となる。 なお、 F D Tは概ね 1000eC程度までとするのが好ましい。 The finish-rolling exit temperature FDT should be 800 ° C or higher in order to obtain a uniform and fine hot-rolled base plate structure. If the FDT is lower than 800 ° C, the structure of the steel sheet becomes non-uniform, a part of the processed structure remains, and after the cold rolling annealing process, the non-uniform structure of the steel remains without disappearing. For this reason, there is a great risk that various problems occur during press forming. In addition, if a high winding temperature is used to avoid the remaining of the processed tissue, coarse crystals are generated, and the same problem occurs. In addition, when the winding temperature is set to a high temperature, the amount of solute N is remarkably reduced, so that it is difficult to obtain a target tensile strength of 440 MPa or more. For this reason, the finish rolling exit temperature FDT was set to 800 ° C or higher. To further improve the mechanical properties, it is desirable that the FDT be 820 ° C or higher. In particular, the upper limit of FDT is not specified, but if it is excessively high, scale flaws and the like will be noticeable. Incidentally, FDT is generally preferably up to about 1000 e C.
なお、 仕上圧延後の冷却は特に厳しく限定しないが、 以下の条件が鋼板の長手 · 幅方向の材質均一性の点で望ましい。 すなわち、 本発明では、 仕上圧延終了後直ち に （0. 5 秒以内に） 冷却を開始し、 冷却中の平均冷却速度を 40 s以上とするの が望ましい。 この条件を満足させることにより、 A1N が析出する高温域を急冷でき、 固溶状態の Nを有効に確保できる。 この冷却開始時間または冷却速度が、 上記条件 を満足しない場合には、粒成長が進みすぎて結晶粒径の微細化が達成しにくいうえ、 圧延で導入された歪エネルギーによる A1N の析出が促進される傾向にあり、 固溶 N 量が欠乏する恐れがあり、 組織が不均一化する傾向となる。 なお、 材質 .形状の均 一性を確保する観点からは、 冷却速度は 300 eCZ s以下に抑えるのが好ましい。 卷取温度： 800 °C以下 The cooling after finish rolling is not particularly strictly limited, but the following conditions are desirable from the viewpoint of the material uniformity in the longitudinal and width directions of the steel sheet. That is, in the present invention, it is desirable that cooling be started immediately (within 0.5 seconds) after finishing rolling and the average cooling rate during cooling be 40 s or more. By satisfying this condition, the high-temperature region where A1N precipitates can be rapidly cooled, and N in solid solution can be secured effectively. If the cooling start time or cooling rate does not satisfy the above conditions, the grain growth will proceed too much, making it difficult to achieve a fine grain size, and the precipitation of A1N due to the strain energy introduced during rolling will be promoted. There is a possibility that the amount of solute N may be deficient, and the tissue tends to be uneven. From the viewpoint of ensuring uniformity of the material and the shape, the cooling rate is preferably suppressed to 300 eCZs or less. Winding temperature: 800 ° C or less
卷取温度 C Tの低下につれて、 鋼板強度が増加する傾向にある。 目標の引張強さ T S 440MPa以上を確保するためには、 C Tは 800 °C以下とするのが好ましい。 なお、 C Tが 200 未満では鋼板形状が乱れやすくなり、 実操業上、 不具合を生じる危険 性が高く、 材質の均一性が低下する傾向を示す。 このため、 C Tは 200で以上とする のが望ましい。 なお、 より材質の均一性が要求される場合には、 C Tは 300 t以上 とするのが好ましい。 なお、 より好ましくは 350 °C以上である。 また、 本発明で は、 仕上圧延において、 熱間圧延荷重を低減するために、 潤滑圧延を行ってもよい。 潤滑圧延を行うことにより、 熱延板の形状 ·材質がより均一化されるという効果が ある。 なお、 潤滑圧延の際の摩擦係数は 0. 25〜0. 10の範囲とするのが好ましい。 ま た、 潤滑圧延と連続圧延とを組み合わせることによりさらに、 熱間圧延の操業が安 定する。  As the winding temperature CT decreases, the steel sheet strength tends to increase. In order to secure the target tensile strength T S 440 MPa or more, it is preferable that C T be 800 ° C. or less. If the CT is less than 200, the shape of the steel sheet is likely to be disturbed, and there is a high risk of causing a problem in actual operation, and the uniformity of the material tends to decrease. Therefore, it is desirable that CT is 200 or more. When more uniformity of the material is required, it is preferable that CT is 300 t or more. The temperature is more preferably 350 ° C. or higher. In the present invention, in finish rolling, lubricating rolling may be performed in order to reduce the hot rolling load. By performing lubricating rolling, there is an effect that the shape and material of the hot rolled sheet are made more uniform. The coefficient of friction during lubrication rolling is preferably in the range of 0.25 to 0.10. In addition, the combination of lubrication rolling and continuous rolling further stabilizes the operation of hot rolling.
上記した熱間圧延工程を施された熱延板は、 ついで、 冷間圧延工程により、 酸洗 およぴ冷間圧延を施されて冷延板となる p  The hot-rolled sheet that has been subjected to the above-mentioned hot rolling step is then subjected to pickling and cold rolling in a cold-rolling step to become a cold-rolled sheet.
酸洗の条件は通常公知の条件でよく、 とくに限定されない。 なお、 熱延板のスケ ールが極めて薄い場合には、 酸洗を施すことなく直ちに冷間圧延を行ってもよい。 また、 冷間圧延条件は、 通常公知の条件でよく、 とくに限定されない。 なお、 組 織の均一性確保という観点から冷間圧下率は 40%以上とするのが好ましい。 つぎ に、 冷間圧延工程の条件限定理由について説明する。  The conditions for pickling may be generally known conditions, and are not particularly limited. If the scale of the hot rolled sheet is extremely thin, cold rolling may be performed immediately without performing pickling. In addition, the cold rolling conditions may be generally known conditions, and are not particularly limited. It is preferable that the cold rolling reduction is 40% or more from the viewpoint of ensuring the uniformity of the tissue. Next, the reasons for limiting the conditions of the cold rolling process will be described.
冷延板は、 ついで箱焼鈍、 連続焼鈍からなる冷延板焼鈍工程を施される。  The cold-rolled sheet is then subjected to a cold-rolled sheet annealing process including box annealing and continuous annealing.
箱焼鈍温度：再結晶温度以上 800 以下  Box annealing temperature: not less than recrystallization temperature and not more than 800
本発明では、 冷延板に箱焼鈍を施し、 素地となるフェライ ト相の集合組織を制御 する。 このフェライ ト相の集合組織制御により製品板の高 r値化が図れる。 この箱 焼鈍により、 製品板には高 r値化に望ましい （1 1 1 ) 集合組織が形成されやすく なる。  In the present invention, box annealing is performed on the cold-rolled sheet to control the texture of the ferrite phase as a base. By controlling the texture of the ferrite phase, the r-value of the product plate can be increased. This box annealing facilitates the formation of a (11 1) texture that is desirable for increasing the r-value on the product sheet.
箱焼鈍温度が再結晶温度未満では、 再結晶が完了せず、 フ ライ ト相の集合組織 を調整することができず、 高 r値化が図れない。 一方、 800 でを超える温度で箱焼 鈍すると、 鋼板の表面欠陥の発生が顕著となり、 初期の目的が達成できなくなる。 なお、 箱焼鈍は、 窒素ガスを主体とし、 3〜5 %の水素ガスを含む焼鈍雰囲気で行 うのが好ましく、 この場合、加熱 ·冷却速度は通常の箱焼鈍の条件でよく、概ね 30°C Ar 程度となる。 また、 焼鈍雰囲気ガスを 100 %水素ガスとすることにより、 より 速い加熱 ·冷却速度としてもよい。 If the box annealing temperature is lower than the recrystallization temperature, recrystallization is not completed and the texture of the fly phase Cannot be adjusted, and a high r-value cannot be achieved. On the other hand, if box annealing is performed at a temperature exceeding 800, the occurrence of surface defects on the steel sheet becomes remarkable, and the initial purpose cannot be achieved. The box annealing is preferably performed in an annealing atmosphere mainly containing nitrogen gas and containing 3 to 5% of hydrogen gas.In this case, the heating and cooling rates may be the same as those in ordinary box annealing, and are generally about 30 °. It is about C Ar. Further, by using an annealing atmosphere gas of 100% hydrogen gas, a higher heating / cooling rate may be obtained.
連続焼鈍温度： Ac 1変態点以上 （Ac 3変態点— 20°C) 以下  Continuous annealing temperature: Ac 1 transformation point or more (Ac 3 transformation point-20 ° C) or less
連続焼鈍温度が Ac 1変態点未満では、 焼鈍後にマルテンサイ ト相が形成されず、 一方、 （Ac 3変態点一 20で） を超えると、 箱焼鈍で形成した望ましい集合組織が変 態により失われるため、 高 r値を有する製品板が得られない。 このため、 連続焼鈍 温度は Ac 1変態点以上 （Ac 3変態点一 20 ） 以下とするのが好ましい。 また、 連続焼鈍時間の保持時間は、 生産効率、 組織の微細化、 固溶 N量の確保の観点から、 できる限り短いほうが好ましい。 一方、 操業の安定性の観点から、 保持時間は 10 s 以上とするのが好ましく、 また、 組織の微細化と固溶 N量の確保という観点からは、 120 s以下とすることが好ましい。 なお、 材質の安定化という観点からは、 20 s以 上とするのがより好ましい。  If the continuous annealing temperature is lower than the Ac 1 transformation point, no martensite phase is formed after annealing, while if it exceeds (at the Ac 3 transformation point-20), the desired texture formed by box annealing is lost by the transformation. Therefore, a product plate having a high r value cannot be obtained. For this reason, it is preferable that the continuous annealing temperature is not less than the Ac 1 transformation point and not more than the Ac 3 transformation point. Further, the holding time of the continuous annealing time is preferably as short as possible from the viewpoints of production efficiency, refining the structure, and securing the amount of solute N. On the other hand, the holding time is preferably 10 s or more from the viewpoint of operation stability, and is preferably 120 s or less from the viewpoint of refining the structure and securing the amount of dissolved N. In addition, from the viewpoint of stabilizing the material, it is more preferable to set the time to 20 s or longer.
連続焼鈍後の冷却： 500 °C以下の温度域まで 10〜 300で の冷却速度で冷却 連 続焼鈍における均熱後の冷却は、 組織の微細化、 マルテンサイ トの形成、 固溶 N量 の確保の観点から重要である。 本発明では、 少なくとも、 500 °C以下の温度域まで 1 O s以上の冷却速度で連続冷却する。 冷却速度が lO Vs未満では、 必要量のマルテ ンサイ ト量と、 均一でかつ微細な組織と、 十分な量の固溶 Nを得ることができない。 一方、 冷却速度が 300 /sを超えると、 過飽和な固溶 C量が顕著に増加するため鋼 板の幅.方向での材質の均一性が低下する。 連続焼鈍後の 10〜 300°C/sの冷却速度で の冷却停止温度が 500 ¾を超えると、 組織の微細化が達成できない。  Cooling after continuous annealing: Cooling down to a temperature range of 500 ° C or less at a cooling rate of 10 to 300 Cooling after soaking in continuous annealing reduces the size of the structure, forms martensite, and secures the amount of solid solution N. It is important from the point of view. In the present invention, continuous cooling is performed at a cooling rate of 1 Os or more to at least a temperature range of 500 ° C or less. If the cooling rate is less than 10 Vs, the required amount of martensite, a uniform and fine structure, and a sufficient amount of solute N cannot be obtained. On the other hand, when the cooling rate exceeds 300 / s, the amount of supersaturated solid solution C increases remarkably, and the uniformity of the material in the width direction of the steel sheet decreases. If the cooling stop temperature at a cooling rate of 10 to 300 ° C / s after continuous annealing exceeds 500 ° C, microstructure refinement cannot be achieved.
過時効処理条件：連続焼鈍後の冷却に引き続き、該冷却の冷却停止温度以下 350 V 以上の温度域で滞留時間 20 s以上 連続焼鈍の均熱後の冷却の冷却停止に引き続き、 冷却停止温度以下 350 °C以上の 温度域で滞留時間 20 s以上の過時効処理を行っても良い。 過時効処理を行うことに より、 固溶 N量を維持したまま、 固溶 C量を選択的に低減することができる。 滞留 温度域が 350 未満では、 固溶 Cの低減に長時間を要し、 生産性低下に繋がるため、 350 °C以上の温度域とするのが好ましい。 Overaging treatment condition: After cooling after continuous annealing, residence time is 20 s or more in the temperature range of 350 V or less, which is lower than the cooling stop temperature of the cooling. Subsequent to the cooling stop after the soaking in the continuous annealing, an overaging treatment with a residence time of 20 s or more may be performed in a temperature range of 350 ° C or lower below the cooling stop temperature. By performing the overaging treatment, the amount of solute C can be selectively reduced while maintaining the amount of solute N. If the residence temperature range is lower than 350, it takes a long time to reduce the solid solution C, which leads to a decrease in productivity. Therefore, the temperature range is preferably 350 ° C or higher.
冷却停止温度以下 350 °C以上の温度域で 20 s以上滞留させることにより、 固溶 C 量を低減でき、 より高度の室温での非時効化が達成される。 滞留時間をより長くす ることにより、 更なる改善が望めるが、 概ね 120 s程度でその効果は飽和する傾向 にあるため、 滞留時間は 120 s以下とするのが好ましい。  By staying for 20 s or more in the temperature range below the cooling stop temperature and 350 ° C or more, the amount of solid solution C can be reduced, and a higher degree of non-aging at room temperature can be achieved. Further improvement can be expected by making the residence time longer, but the effect tends to be saturated at about 120 s. Therefore, the residence time is preferably 120 s or less.
大きな歪時効硬化量を得るためには、 固溶 Cと固溶 Nのどちらも利用することが 有利であるが、 固溶 Cを利用すると、 室温での時効劣化が顕著となり、 鋼板の適用 部位が制限されることになる。 従って、 汎用性のある歪時効硬化型鋼板を製造する には、 充分な量の固溶 Nを確保した上で過時効処理を行うのが好ましい。  In order to obtain a large amount of strain age hardening, it is advantageous to use both solid solution C and solid solution N.However, when solid solution C is used, aging deterioration at room temperature becomes remarkable, Will be limited. Therefore, in order to produce a versatile strain-age hardened steel sheet, it is preferable to perform overaging treatment after securing a sufficient amount of solid solution N.
なお、 本発明の高張力冷延鋼板の表面に溶融めつき層を有する高張力冷延めっき 鋼板を製造する場合、 箱焼鈍についで行う連続焼鈍を連続溶融めつきラインにて行 い、 連続焼鈍後の冷却に引き続いて溶融亜鉛めつき、 あるいはさらに合金化処理を 行い、 溶融亜鉛めつき鋼板を製造することもできる。  When manufacturing a high-tensile cold-rolled steel sheet having a hot-dip layer on the surface of the high-tensile cold-rolled steel sheet of the present invention, continuous annealing following box annealing is performed in a continuous hot-dip line, and continuous annealing is performed. Subsequent to subsequent cooling, hot-dip galvanizing or further alloying can be performed to produce a hot-dip galvanized steel sheet.
調質圧延あるいはレべラー加工：伸び率 0. 2 〜15%  Temper rolling or leveling: elongation 0.2 to 15%
本発明では、 冷延焼鈍工程に引き続いて、 形状矯正、 粗度調整の目的で、 調質圧 延またはレべラー加工を施してもよい。 調質圧延あるいはレべラー加工の伸び率が 合計で 0. 2 %未満では、 形状矯正、 粗度調整の所期の目的が達成できない。 一方、 1 5%を超えると、 顕著な延性の低下をもたらす。 なお、 調質圧延とレべラー加工では その加工形式が相違するが、 その効果は両者で大きな差異がないことを確認してい る。 調質圧延、 レべラー加工は、 めっき処理後でも有効である。 以下， 参考のため、 この発明鋼板をプレス成形などの成形加工に供した場合にお ける成形条件おょぴその後の強度上昇熱処理条件について説明する。 この発明の 鋼板を、 例えば絞り加工などのプレス加工に供する場合、 プレス加工により導入さ れる歪みは数％〜十数％である。 成形部品によって歪み量は変化するが、 自動車分 野における内板および構造部材は 5〜10%程度の歪みが導入される。 In the present invention, after the cold rolling annealing step, temper rolling or leveling may be performed for the purpose of shape correction and roughness adjustment. If the total elongation of the temper rolling or leveling is less than 0.2%, the intended purposes of shape correction and roughness adjustment cannot be achieved. On the other hand, if it exceeds 15%, the ductility is significantly reduced. Although the form of temper rolling and leveler processing are different, it has been confirmed that there is no significant difference between the two. Temper rolling and leveling are effective even after plating. Hereinafter, for reference, when the steel sheet of the present invention is subjected to forming such as press forming, The molding conditions and the subsequent heat treatment conditions for increasing the strength will be described. When the steel sheet of the present invention is subjected to press working such as drawing, for example, the strain introduced by press working is several percent to several tens of percent. Although the amount of distortion varies depending on the molded parts, about 5 to 10% of distortion is introduced into the inner plates and structural members in the automotive field.
ついで、 これらの成形部品には、 塗装焼付け処理などの熱処理が施されるが、 こ の発明鋼板では熱処理後に成形品強度を効果的に高めることができる。 なお、 こ の発明では、 かような焼付硬化性を実験室にて評価する方法として、 JIS 5号サイ ズの引張試験片を圧延方向に採取し、 引張試験機により 10%の引張歪を付与し、 そ の後熱処理を加えたのち、 再度引張り試験を実施する。 特に低温域での熱処理後の 特性を評価する場合は、 熱処理条件を 120 ， 20分とする。 この試験は、 プレス成 形に引き続き熱処理を行った完成後の部位の特性を評価するものである。  Then, these formed parts are subjected to a heat treatment such as a paint baking treatment. With the steel sheet of the present invention, the strength of the formed parts can be effectively increased after the heat treatment. In this invention, as a method of evaluating such bake hardenability in a laboratory, a tensile test piece of JIS No. 5 size was sampled in the rolling direction, and 10% tensile strain was applied by a tensile tester. After that, heat treatment is performed, and then the tensile test is performed again. In particular, when evaluating the characteristics after heat treatment in the low temperature range, heat treatment conditions should be 120 and 20 minutes. This test evaluates the properties of the completed parts that have been subjected to heat treatment following press forming.
すなわち、 この発明では、 このような引張り歪付与一熱処理後の引張強度と製品 の引張強度との差 （A TS) を強度上昇熱処理能として定義する。  That is, in the present invention, the difference (ATS) between the tensile strength after such a tensile strain imparting-heat treatment and the tensile strength of the product is defined as the strength increasing heat treatment ability.
通常、 成形品の強度上昇を高めるには、 成形により導入する歪み量が大きいまた は加工後の熱処理温度が高い方が好ましい。  Usually, in order to increase the strength of the molded product, it is preferable that the amount of distortion introduced by molding is large or the heat treatment temperature after working is high.
しかしながら、 この発明鋼板は、 付与歪み量が上記した 5〜10%程度の場合に、 従来よりも成形後熱処理温度が低くても、 すなわち熱処理温度が 200°C以下であつ ても、 十分な強度の上昇を図ることができる。 とはいえ、 熱処理温度が 120°C未満 では歪みが低い場合に十分な強度上昇効果が得られない。 一方、 成形後の熱処理温 度が 350eCを超える温度になると軟化が進行する。 従って、 成形後の熱処理温度は 120-350 程度とするのが好ましい。 However, the steel sheet of the present invention has sufficient strength even when the heat treatment temperature after forming is lower than before, that is, even when the heat treatment temperature is 200 ° C or less, when the applied strain amount is about 5 to 10% described above. Can be increased. Nevertheless, if the heat treatment temperature is lower than 120 ° C, a sufficient strength increasing effect cannot be obtained when the strain is low. On the other hand, when the heat treatment temperature after molding exceeds 350 eC , softening proceeds. Therefore, it is preferable that the heat treatment temperature after molding be about 120-350.
なお、 加熱方法としては、 熱風加熱、 赤外炉加熱、 温浴熱処理、 通電加熱、 高周 波加熱などの方法が適用でき、 特に規定されない。 また、 強度を上昇させたい部分 のみを選択的に加熱する場合でもよい。 実施例 以下の実施例において、 固溶 N量、 微視組織、 引張特性、 r値測定、 歪時効硬化 特性、 時効特性を調査した。 調査方法は下記のとおりである。 As a heating method, a method such as hot air heating, infrared furnace heating, hot bath heat treatment, electric current heating, and high frequency heating can be applied, and is not particularly specified. Alternatively, only the portion where the strength is to be increased may be selectively heated. Example In the following examples, the amount of solid solution N, microstructure, tensile properties, r value measurement, strain age hardening properties, and aging properties were investigated. The survey method is as follows.
(1) 固溶 N量  (1) Solid solution N content
固溶 N量は、 化学分析により求めた鋼中の全 N量から析出 N量を差し引いて求め た。 ここで析出 N量は、 上記した定電位電解法を用いた分析法により求めた。  The amount of solute N was determined by subtracting the amount of precipitated N from the total amount of N in the steel determined by chemical analysis. Here, the amount of precipitated N was determined by an analytical method using the above-described potentiostatic electrolysis method.
(2) 微視組織  (2) Microstructure
各冷延焼鈍板から試験片を採取し、 圧延方向に直交する断面 （C断面） について、 光学顕微鏡あるいは走査型電子顕微鏡を用いて微視組織を撮像し、 画像解析装置を 用いてフェライトの組織分率および第 2相の種類およぴ組織分率を求めた。  A specimen was taken from each cold-rolled annealed plate, and the microstructure of the cross section (C cross section) perpendicular to the rolling direction was imaged using an optical microscope or a scanning electron microscope, and the microstructure of ferrite was obtained using an image analyzer. The fraction, the type of the second phase and the tissue fraction were determined.
(3) 結晶粒径  (3) Grain size
本発明では結晶粒径として、 断面組織写真から AS TMに規定の求積法により算 出した値と、断面組織写真から AS TMに規定の切断方法により求めた公称粒径（例 えば梅本ら ：熱処理、 24 (1 9 8 4) 、 3 34参照） のうち、 いずれか大きい方 を採用した。  In the present invention, as the crystal grain size, a value calculated by a quadrature method specified in ASTM from a cross-sectional structure photograph and a nominal particle size determined by a cutting method specified in ASTM from a cross-sectional structure photograph (for example, Umemoto et al .: Heat treatment, 24 (1994), 334) was used, whichever was greater.
(4) 引張特性  (4) Tensile properties
各冷延焼鈍板から JIS 5号試験片を圧延方向に採取し、 JIS Z 2241の規定に準拠 して歪速度： 3 Xl0_ 3/sで引張試験を実施し、 降伏応力 YS、 引張強さ T S、 伸び E 1を求めた。 The JIS 5 test piece No. taken in the rolling direction from each cold rolled annealed sheet, the distortion in compliance with the provisions of JIS Z 2241 Speed: 3 to a tensile test at XL0 _ 3 / s, the yield stress YS, tensile strength TS and elongation E1 were determined.
(5) 歪時効硬化特性  (5) Strain age hardening characteristics
各冷延焼鈍板から JIS 5号試験片を圧延方向に採取し、 予変形としてここでは 5%の引張予歪を与えて、 ついで 170 °CX20min の塗装焼付処理相当の熱処理を施 したのち、歪速度： 3 X 10— 3 で引張試験を実施し、予変形—塗装焼付処理後の引 張特性（降伏応力 Y S BH、引張強さ TS BH)を求め、 BH量 = YS BH— Y S 5%、 厶 TS=T S BH— TSを算出した。 なお、 YS 5%は、製品板を 5 %予変形したと きの変形応力であり、 YS BH、 T SBHは予変形一塗装焼付処理後の降伏応力、 引 張強さであり、 T Sは製品板の引張強さである。 (6) r値測定 A JIS No. 5 test piece was sampled from each cold-rolled annealed sheet in the rolling direction, a 5% tensile pre-strain was given here as a pre-deformation, and then a heat treatment equivalent to paint baking at 170 ° C for 20 minutes was performed, followed by distortion. speed: 3 to a tensile test at X 10- 3, predeforming - tensile properties (yield stress YS BH, tensile strength TS BH) after paint baking look, BH amount = YS BH- YS 5%, TS = TS BH—TS was calculated. YS 5% is the deformation stress when the product plate is pre-deformed 5%, YS BH and T SBH are the yield stress and tensile strength after pre-deformation-paint baking, and TS is the product plate Is the tensile strength. (6) r value measurement
各冷延焼鈍板の圧延方向 （L方向） 、 圧延方向に対し 45° 方向 （D方向） 、 圧延 方向に対し 90° 方向 （C方向） から、 JIS 5 号試験片を採取した。 これら試験片に 1 5%の単軸引張予歪を付与した時の各試験片の幅歪と板厚歪を求め、 r値の定義式で ある、 幅歪と板厚歪の比、  JIS No. 5 test pieces were sampled from the rolling direction (L direction), 45 ° direction (D direction) with respect to the rolling direction, and 90 ° direction (C direction) with respect to the rolling direction of each cold rolled annealed sheet. When 15% uniaxial tensile prestrain is applied to these test pieces, the width strain and the thickness strain of each test piece are obtained, and the ratio of the width strain to the thickness strain, which is the definition formula of r value,
r =ln (w/w 0 ) /In ( t / t 0 )  r = ln (w / w 0) / In (t / t 0)
(ここで、 wo 、 t O は試験前の試験片の幅おょぴ板厚であり、 w、. tは試験後 の試験片の幅おょぴ板厚である。 ）  (Here, wo and t O are the width and thickness of the test piece before the test, and w and .t are the width and thickness of the test piece after the test.)
から各方向の r値を求め、 次式 Calculate the r value in each direction from
r me a n= (T L +2 r D + r c . ) Z4 r me an = (TL +2 r D + r c .) Z4
により平均 r値 r m e a nを求めた。 ここで、 r L は、 圧延方向 （L方向） の r値 であり、 rD は、 圧延方向 （L方向） に対し 45° 方向 （D方向） の r値であり、 r c は、 圧延方向 （L方向） に対し 90° 方向 （C方向） の r値である。 なお、 実験の 精度の向上のため、 体積一定と仮定して、 伸ぴ歪と幅方向の歪の変化で算出した。The mean r value r mean was determined by Here, r L is the r value in the rolling direction (L direction), r D is the r value in the 45 ° direction (D direction) with respect to the rolling direction (L direction), and rc is the rolling direction ( R value in the 90 ° direction (C direction) with respect to the L direction). In addition, in order to improve the precision of the experiment, assuming that the volume was constant, the calculation was performed based on changes in the elongation strain and the strain in the width direction.
(7) 時効特性 (7) Aging characteristics
各冷延焼鈍板から JIS 5 号試験片を採取し、 該試験片に、 50^X2001^ の時効処 理を施したのち、 引張試験を実施した。 得られた結果から、 時効処理前後の降伏伸 ぴ差 ΔΥ-Elを求め、 常温の時効特性を評価した。 ΔΥ- E1が零であれば非時効性とし て耐常温時効特性に優れると評価した。  JIS No. 5 test pieces were collected from each cold-rolled annealed sheet, and the test pieces were subjected to an aging treatment of 50 ^ X2001 ^ and then subjected to a tensile test. From the obtained results, the yield-elongation difference ΔΥ-El before and after the aging treatment was determined, and the aging characteristics at room temperature were evaluated. If ΔΥ-E1 was zero, it was evaluated as non-aging and excellent in normal temperature aging resistance.
(8) 成形一熱処理後の引張強度  (8) Tensile strength after molding and heat treatment
成形一熱処理後の引張強度は、 製品板から圧延方向に JIS 5 号試験片を採取し、 予歪み 10%を付与した後、 120°Cおよび従来から行われている塗装焼付相当熱処理 温度である 170°Cにて 20分間の熱処理を施し、 引張強度を測定して求めた。  The tensile strength after forming and heat treatment is the temperature at 120 ° C and the heat treatment temperature equivalent to conventional paint baking after taking a JIS No. 5 test piece from the product plate in the rolling direction and applying a 10% pre-strain. Heat treatment was performed at 170 ° C for 20 minutes, and the tensile strength was measured and determined.
(9) 常温時効による全伸びの低下量 （ΔΕ1)  (9) Decrease in total elongation due to normal temperature aging (ΔΕ1)
常温時効による全伸びの低下量 （ΔΕ1) は、 製品板から圧延方向に JIS 5 号試験 片を採取して測定した全伸びと、 別途、 圧延方向に採取した JIS5 号試験片を用い常 温時効の促進処理 （ 100°C， 8時間保持） を施したのちに測定した全伸びとの差と して求めた。 The amount of decrease in total elongation due to normal temperature aging (Δ 効 1) was calculated using the total elongation measured by taking a JIS No. 5 test piece from the product plate in the rolling direction and the JIS No. 5 test piece separately taken in the rolling direction. The difference from the total elongation measured after accelerating the heat aging treatment (holding at 100 ° C for 8 hours) was obtained.
実施例 1 Example 1
表 1に示す成分組成になる鋼スラブを、 表 2に示す条件で板厚： 3. 5 匪の熱延板、 ついで板厚： 0. 7 rainの冷延板としたのち、 連続焼鈍ラインまたは連続焼鈍一合金化 溶融亜鉛めつきラインにて再結晶焼鈍、 さらには合金化溶融亜鉛めつき処理を施し、 その後圧下率： 1. 0 %の調質圧延を施して、 冷延鋼板および片面当たりの目付量： 4 A steel slab having the composition shown in Table 1 was converted into a hot-rolled strip with a thickness of 3.5 and then a cold-rolled strip with a thickness of 0.7 rain under the conditions shown in Table 2. Continuous annealing-alloying Recrystallization annealing at the hot-dip galvanizing line, further hot-dip galvanizing treatment, then temper rolling at a rolling reduction of 1.0%, per cold rolled steel sheet and one side Weight per unit area: 4
5g/m 2で両面めつきした合金化溶融亜鉛めつき鋼板を製造した。 なお、 表 2のうち N o. 3, 8 の熱延仕上終了温度は Ar 3変態点未満であり、 それ以外は Ar s変態点以上 である。 An alloyed hot-dip galvanized steel sheet plated on both sides at 5 g / m 2 was produced. In Table 2, the finish temperatures of hot rolling of Nos. 3 and 8 are lower than the Ar 3 transformation point, and the other temperatures are higher than the Ars transformation point.
かく して得られた冷延鋼板おょぴ合金化溶融亜鉛めつき鋼板の引張強度おょぴ r 値、 ならびに成形一熱処理後の引張強度の変化について調査した結果を、 表 3に示 す。  Table 3 shows the results of an investigation on the tensile strength and the r-value of the cold-rolled steel sheet and the alloyed hot-dip galvanized steel sheet thus obtained, and the change in the tensile strength after the forming and heat treatment.
表 3から明らかなように、 この発明に従い得られた冷延鋼板および合金化溶融鉛 めっき鋼板はいずれも、 比較例に比べて、 高い r値と優れた歪み時効硬化特性が得 られている。 また、 特に適合例のうち、 結晶粒径が 20 /z ra 以下のものは、 常温時効 による伸びの低下量も Δ Ε1で 2. 0%以下と小さくなっている。  As is evident from Table 3, both the cold-rolled steel sheet and the alloyed hot-dip galvanized steel sheet obtained according to the present invention have higher r-values and better strain age hardening characteristics than the comparative examples. Particularly, among the conforming examples, those having a crystal grain size of 20 / z ra or less have a small decrease in elongation due to aging at room temperature of 2.0% or less at ΔΕ1.
実施例 2 Example 2
表 1に記載の鋼記号 Bのスラブを用い、 表 2の No. 2と同じ製造条件であるスラブ 加熱温度： 1100で、 仕上熱延温度： 900 でで熱延したのち、 卷取り温度： 550 でコイルに卷き取った。 このコイルを、 圧下率： 80%で冷間圧延した後、 840 で再結晶焼鈍を行った。 得られた冷延鋼板の製品特性は、 引張強度 TS = 365 MPa 、 r値 = 1. 7 であった。 この冷延鋼板から JIS 5号試験片を圧延方向に採取し、 引 張試験機により 10%の引張歪みを付与したのち、 表 4に示す熱処理条件 （温度、 時 間） で熱処理を実施し、 再度引張り試験を行った。 表 4に、 歪付与前の製品の引 張強度 （TS = 365 MPa)からの引張強度の上昇代 （A TS) を併記する。 Using a slab of steel symbol B shown in Table 1, the slab is manufactured under the same manufacturing conditions as No. 2 in Table 2, heating temperature: 1100, finishing hot rolling temperature: 900, and rolling temperature: 550 And wound on the coil. The coil was cold-rolled at a rolling reduction of 80%, and then recrystallized at 840. The product properties of the obtained cold-rolled steel sheet were as follows: tensile strength TS = 365 MPa, r value = 1.7. A JIS No. 5 test piece was sampled from the cold-rolled steel sheet in the rolling direction, subjected to a 10% tensile strain by a tensile tester, and then subjected to heat treatment under the heat treatment conditions (temperature, time) shown in Table 4. The tensile test was performed again. Table 4 shows the product drawing before strain application. The rise in tensile strength (ATS) from the tensile strength (TS = 365 MPa) is also shown.
表 4に示したとおり、 強度の上昇量は、 熱処理温度が高くなるほど、 また熱処理 時間が長くなるほど大きくなるが、 発明鋼板は熱処理温度が 120°Cと低温で、 かつ 保持時間が 2分と短くても 82 MPa という十分な引張強度の上昇 （20分熱処理時の 8 5%以上） が得られ、 低温，短時間の熱処理でも良好な歪み時効硬化特性が得られる ことが分かる。 なお、 自動車の構造部材等において、 安定した強度上昇効果を得 るために、 通常の温度、 時間で熱処理を行うことに何ら間題はない。 また、 この 冷延鋼板に対して溶融亜鉛めつきおょぴ加熱合金化処理を施して得た合金化溶融亜 鉛めつき鋼板についても、 表 4と同様な結果が得られることが確かめられている。 実施例 3  As shown in Table 4, the increase in strength increases as the heat treatment temperature increases and as the heat treatment time increases, but the invention steel sheet has a low heat treatment temperature of 120 ° C and a short holding time of 2 minutes. However, a sufficient increase in tensile strength of 82 MPa (more than 85% during 20-minute heat treatment) was obtained, indicating that good strain age hardening characteristics can be obtained even at low temperatures and short-time heat treatment. It should be noted that there is no problem in performing heat treatment at normal temperature and time in order to obtain a stable strength increase effect for structural members of automobiles. It was also confirmed that the same results as in Table 4 were obtained for the alloyed molten galvanized steel sheet obtained by subjecting this cold-rolled steel sheet to a hot-dip galvanizing heat alloying treatment. I have. Example 3
表 6に示す組成になる鋼スラブを表 7に示す条件で熱間圧延し板厚 3. 5mra の熱延 板とした。 これらの熱延板を表 7に示す条件で冷間圧延して板厚 0. 7mm の冷延板と し、 これらの冷延板を同表に示す条件で再結晶焼鈍し、 うち一部につきさらに同表 に示す条件で溶融亜鉛めつきあるいは合金化溶融亜鉛めつきを施した。 得られた 製品板について、 固溶 N量、 微視組織、 引張特性、 歪時効硬化特性を調査した。 結果を表 8に示す。 同表より、 本発明による鋼板は、 T S X r値≥750MPa ( Bと N b, Ti，V の 1種または 2種以上とを複合添加したものではさらに T S X r値≥850MP a) 、 B H≥80MPa 、 A T S≥40MPa をすぺて満たすが、 比較例ではこれら 3特性の 1つ以上が本発明のレベルに達しない。  A steel slab having the composition shown in Table 6 was hot-rolled under the conditions shown in Table 7 to obtain a hot-rolled sheet with a thickness of 3.5 mra. These hot-rolled sheets were cold-rolled under the conditions shown in Table 7 to obtain 0.7 mm-thick cold-rolled sheets, and these cold-rolled sheets were recrystallized and annealed under the conditions shown in the same table. Further, hot-dip galvanizing or alloyed hot-dip galvanizing was performed under the conditions shown in the same table. The obtained product plate was examined for the amount of solute N, microstructure, tensile properties, and strain age hardening properties. Table 8 shows the results. According to the table, the steel sheet according to the present invention has a TSX r value of ≥750 MPa (for those in which B and one or more of Nb, Ti, and V are added in combination, a TSX r value of ≥850 MPa) and a BH≥80 MPa ATS ≥40 MPa, but one or more of these three characteristics do not reach the level of the present invention in the comparative example.
実施例 4 Example 4
次に本発明の実施例について説明する。  Next, examples of the present invention will be described.
表 9に示す組成の溶鋼を転炉で溶製し、 連続铸造法で鋼スラブとした。 これら鋼 スラブを表 1 0に示す条件で加熱し、 粗圧延してシートパーとし、 ついで表 1 0に 示す条件の仕上げ圧延を施す熱間圧延工程により熱延板とした。 なお、 Ar 3変態点 を熱間仕上げ圧延条件をシミュレートした条件で加工変態測定装置 （富士電波ェ機 製） を用いて測定し、 表 1 0に示した。 これら熱延板を酸洗および表 1 0に示す条件の冷間圧延からなる冷間圧延工程に より冷延板とした。 ついで、 これら冷延板に表 1 0に示す条件で連続焼鈍を行った。 一部については、 冷延焼鈍工程につづいて、 調質圧延を施した。 Molten steel with the composition shown in Table 9 was smelted in a converter and made into a steel slab by continuous casting. These steel slabs were heated under the conditions shown in Table 10, rough-rolled to form sheet sheets, and then hot-rolled by a hot rolling step of finish rolling under the conditions shown in Table 10. The Ar 3 transformation point was measured using a working transformation measuring device (manufactured by Fuji Denki Co., Ltd.) under conditions simulating the hot finish rolling conditions, and is shown in Table 10. These hot-rolled sheets were formed into cold-rolled sheets by a cold rolling process including pickling and cold rolling under the conditions shown in Table 10. Subsequently, continuous annealing was performed on these cold-rolled sheets under the conditions shown in Table 10 below. For some, temper rolling was performed after the cold rolling annealing process.
得られた冷延焼鈍板について、 固溶 N量、 微視組織、 引張特性、 r値測定、 歪時 効硬化特性、 時効特性を調査した。  With respect to the obtained cold rolled annealed sheet, the amount of solute N, microstructure, tensile properties, r value measurement, strain age hardening properties, and aging properties were investigated.
なお、 No. 4、 No. 10 の鋼板表面に、 溶融亜鉛めつきを施しめつき鋼板とし、 同様 に各種特性を評価した。  The surfaces of No. 4 and No. 10 steel sheets were subjected to hot-dip galvanization, and various properties were evaluated in the same manner.
これらの結果を表 1 1に示す。  Table 11 shows the results.
本発明例では、 いずれも優れた延性と、 格段に高い B H量、 A T Sを呈し、 優れ た歪時効硬化特性と、 平均 r値 1. 2 以上の高い r値と、 常温時効では非時効性であ り優れた耐常温時効性を有している。 なお、 No. 4、 No. 10 の鋼板表面に、 溶融亜鉛 めっきを施しためっき鋼板の特性は、 めっき層の幅縮み拘束のため、 冷延鋼板に比 ぺ平均 r値で 0. 2 、 伸ぴ E 1で 1 %程度の特性の低下があつたが、 歪時効硬化性、 耐常温時効性はめつき前の特性と殆ど変化はなかった。 これに対し、 本発明の範囲 を外れる比較例は、 延性が劣化しているか、 B H量、 A T Sが少ないか、 時効劣化 が著しいかで、 目標の特性を全て具備することはなく、 十分な特性を有する鋼板と はいえない。  All of the examples of the present invention exhibit excellent ductility, a remarkably high BH content and ATS, excellent strain aging hardening properties, a high r-value with an average r-value of 1.2 or more, and non-aging properties at room temperature aging. It has excellent room temperature aging resistance. The properties of galvanized steel sheets with hot-dip galvanized No. 4 and No. 10 steel sheets were 0.2 times the average r-value compared to cold-rolled steel sheets and 0.2% due to the restraint of width reduction of the coating layer.ぴ E 1 decreased the properties by about 1%, but the strain age hardening property and the room temperature aging resistance were almost the same as the properties before plating. On the other hand, the comparative examples out of the scope of the present invention do not have all of the target properties and have sufficient properties, depending on whether the ductility is deteriorated, the amount of BH, the ATS is small, or the aging deterioration is remarkable. It cannot be said that the steel sheet has
鋼板 No. 11 は、 C、 Al、 N、 NZA1が本発明範囲を外れ、 そのため、 r値、 B H 量、 A T S、 耐常温時効性が低下している。 また、 鋼板 No. 12 は、 B、 Nbが本発明 範囲を外れ、 ァシキユラ一フェライ ト量が本発明範囲を低くはずれ、 そのため、 B H量、 A T S、 耐常温時効性が低下している。  In the steel sheet No. 11, C, Al, N, and NZA1 were out of the range of the present invention, and therefore, the r value, the BH amount, the ATS, and the aging resistance at room temperature were lowered. In steel sheet No. 12, B and Nb were out of the range of the present invention, and the amount of ferrite was out of the range of the present invention. Therefore, the BH amount, ATS, and the aging resistance at room temperature were lowered.
鋼板 No. 13 は、 Bが本発明の好適範囲を外れ、 ァシキユラ一フェライ ト量が本発 明範囲を低くはずれ、 そのため、 r値、 B H量、 A T S、 耐常温時効性が低下して いる。 また、 鋼板 No. 14 は、 Nbが本発明範囲を外れ、 固溶 N量が本発明範囲を低く はずれ、 そのため、 歪時効硬化特性が低下している。  In the steel sheet No. 13, B was out of the preferred range of the present invention, and the amount of ferrite was out of the range of the present invention. Therefore, the r value, the BH amount, the ATS, and the resistance to room temperature aging were lowered. Further, in steel sheet No. 14, Nb was out of the range of the present invention, and the amount of solute N deviated from the range of the present invention. Therefore, the strain age hardening characteristics were deteriorated.
鋼板 No. 15 は、 Nが本発明の好適範囲を外れ、 固溶 Nが少なく、 歪時効硬化特性 が低下している。 鋼板 No. 17 -No. 20 は、 熱延条件、 冷延板焼鈍条件が好適範囲か ら外れ、 微視組織が本発明の範囲外となり ·Β Η量、 A T Sが低減し、 歪時効硬化特 性が低下し、 耐常温時効性が劣化している。 In steel sheet No. 15, N was out of the preferred range of the present invention, solute N was small, and strain age hardening characteristics Is declining. In the steel sheets No. 17 to No. 20, the hot rolling conditions and the cold rolling annealing conditions were out of the preferred ranges, and the microstructure was out of the range of the present invention. The aging resistance has deteriorated.
実施例 5 Example 5
表 1 2に示す組成になる鋼を、 実施例 4と同様の方法でスラブとなし、 該スラブ を表 1 3に示す条件で加熱し、 粗圧延して 25龍厚のシ^ "トパ一とし、 ついで表 1 3 に示す条件の仕上圧延を施す熱間圧延工程により熱延板とした。 なお、 粗圧延後で 仕上圧延入側で相前後するシートパー同士を溶融圧接法で接合して連続圧延した。 また、 シートパーの幅端部、 長さ方向端部を誘導加熱方式のシートパーエッジヒー タ、 シートパーヒータを使用してシートパーの温度を調節した。  A steel having the composition shown in Table 12 was formed into a slab in the same manner as in Example 4, and the slab was heated under the conditions shown in Table 13 and roughly rolled to obtain a 25-long steel sheet. Then, the hot rolled sheet was formed by a hot rolling process in which finish rolling was performed under the conditions shown in Table 13. After rough rolling, the successive sheet pars on the entry side to the finish rolling were joined by a fusion welding method and continuously rolled. In addition, the width of the sheet par and the end in the length direction of the sheet par were controlled by using an induction heating type sheet par edge heater and a sheet par heater.
これら熱延板を酸洗おょぴ表 1 3に示す条件の冷間圧延からなる冷間圧延工程に より 1. 6 mm厚の冷延板とした。 ついで、 これら冷延板に表 1 3に示す条件で連続焼 鈍を行った。  These hot-rolled sheets were pickled by a cold rolling process comprising cold rolling under the conditions shown in Table 13 to form cold-rolled sheets having a thickness of 1.6 mm. Next, these cold-rolled sheets were subjected to continuous annealing under the conditions shown in Table 13.
得られた冷延焼鈍板について、 実施例 4と同様に固溶 N量、 微視組織、 引張特性、 r値測定、 歪時効硬化特性を調査した。 また、 各冷延焼鈍板の幅方向および長手方 向について、 各 10箇所で引張特性を調査し、 降伏強さ、 引張強 、 伸びのばらつき を調査した。  With respect to the obtained cold rolled annealed sheet, the amount of dissolved N, microstructure, tensile properties, r value measurement, and strain age hardening properties were examined in the same manner as in Example 4. In addition, the tensile properties of each cold-rolled annealed sheet were examined at ten locations in the width direction and the longitudinal direction, and variations in yield strength, tensile strength, and elongation were examined.
それらの結果を表 1 4に示す。  Table 14 shows the results.
本発明例は、 いずれも優れた歪時効硬化特性と高い r値を有し、 製造条件の変動 にもかかわらず、 安定して格段に高い B H量、 A T S、 平均 r値を示した。 また、 本発明例では、 連続圧延とシートパーの長手方向、 幅方向温度調整を実施すること により、 製品鋼板の板厚精度および形状が向上し、 材質ばらつきが 1 / 2に減少す ることを確認した。 また、 調質圧延の伸び率を 0. 5 ~ 2 %、 レべラーの伸び率を 0 〜 1 %まで変化させたが、 歪時効硬化特性の低下はなかった。  Each of the examples of the present invention had excellent strain age hardening characteristics and a high r-value, and showed a stably remarkably high BH content, ATS, and an average r-value despite fluctuations in production conditions. In addition, in the example of the present invention, it was confirmed that the thickness accuracy and the shape of the product steel sheet were improved by performing continuous rolling and temperature adjustment in the longitudinal direction and the width direction of the sheet par, and the material variation was reduced to 1/2. did. The elongation of the temper rolling was changed from 0.5 to 2% and the elongation of the leveler was changed from 0 to 1%, but there was no decrease in strain age hardening characteristics.
実施例 6 Example 6
次に本発明の実施例について説明する。 表 1 5に示す組成の溶鋼を転炉で溶製し、 連続鍀造法で鋼スラブとした。 これら 鋼スラブを表 1 6に示す条件で加熱 （一部温片装入あり） し、 粗圧延してシートパ 一とし、 ついで表 1 6に示す条件の仕上げ圧延を施す熱間圧延工程により熱延板と した。 なお、 一部のシートパーでは相前後するシートパー同士を溶融圧接法で接合 して連続圧延をおこなった。 Next, examples of the present invention will be described. Molten steel with the composition shown in Table 15 was smelted in a converter, and was made into a steel slab by continuous casting. These steel slabs were heated under the conditions shown in Table 16 (with some hot flakes charged), rough-rolled to form a sheet pallet, and then subjected to hot rolling in a hot rolling process in which finish rolling was performed under the conditions shown in Table 16 It was a plate. In addition, in some sheet pars, successive sheet pars were joined to each other by a melt pressure welding method, and continuous rolling was performed.
これら熱延板を酸洗おょぴ表 1 6に示す条件の冷間圧延からなる冷間圧延工程に より冷延板とした。 ついで、 これら冷延板に表 1 6に示す条件で箱焼鈍とそれに続 く連続焼鈍を行った。 一部については、 冷延焼鈍工程につづいて、 調質圧延を施し た。 なお、 箱焼鈍なしの場合も実施した。 箱焼鈍の焼鈍温度は全て再結晶温度以上 とした。 ·  These hot-rolled sheets were pickled by a cold rolling process comprising cold rolling under the conditions shown in Table 16. Next, box annealing and subsequent continuous annealing were performed on these cold-rolled sheets under the conditions shown in Table 16. In some cases, temper rolling was performed after the cold rolling annealing process. The test was also performed without box annealing. The annealing temperatures in box annealing were all higher than the recrystallization temperature. ·
得られた冷延焼鈍板について、 固溶 N量、 微視組織、 引張特性、 r値測定、 歪時 効硬化特性、 時効特性を調査した。  With respect to the obtained cold rolled annealed sheet, the amount of solute N, microstructure, tensile properties, r value measurement, strain age hardening properties, and aging properties were investigated.
なお、 No. 17 、 No. 18 の鋼板表面には、 表中の連続焼鈍の後にインラインで溶融 亜鉛めつきを施しめつき鋼板とし、 同様に各種特性を評価した。  The surfaces of No. 17 and No. 18 steel sheets were subjected to in-line hot-dip galvanizing after continuous annealing as shown in the table, and the various properties were similarly evaluated.
これらの結果を表 1 7に示す。  Table 17 shows these results.
本発明例では、 いずれも優れた延性と、 格段に高い B H量、 A T Sを呈し、 優れ た歪時効硬化特性と、 平均 r値 1. 2 以上の高い r値と、 常温非時効性とを有してい る。 なお、 表 1 7に示す鋼板 No. 17 、 No. 18 の溶融亜鉛めつき鋼板の特性は、 同様 に連続焼鈍した冷延鋼板の特性と殆ど差はなかった。 これに対し、 本発明の範囲を 外れる比較例は、 延性が劣化しているか、 B H量、 A T Sが少ないか、 時効劣化が 著しいかで、 目標の特性を全て具備することはなく、 十分な特性を有する鋼板とは いえない。  All of the examples of the present invention exhibit excellent ductility, a remarkably high BH content and ATS, excellent strain aging hardening properties, a high r value of an average r value of 1.2 or more, and non-aging at room temperature. are doing. The properties of the hot-dip galvanized steel sheets No. 17 and No. 18 shown in Table 17 were almost the same as those of the continuously annealed cold-rolled steel sheets. On the other hand, the comparative examples out of the scope of the present invention do not have all of the target characteristics and have sufficient characteristics depending on whether the ductility is deteriorated, the BH amount and the ATS are small, or the aging deterioration is remarkable. It cannot be said that the steel sheet has
鋼板 No. l l は、 C量、 N量が本発明範囲を外れ、 固溶 N量、 マルテンサイ ト量が 本発明範囲を低くはずれ、 そのため、 B H量、 Δ T Sが低下し Δ Υ - E1が増加してい る。 また、 鋼板 No. 12 は、 Al、 NZA1、 Nが本発明範囲を外れ、 固溶 N量が本発明 範囲を低くはずれ、 フェライ トの平均結晶粒径が本発明の範囲を高く外れ、 そのた め、 B H量、 Δ T Sが低下し Δ Υ - Elが増加している。 In the steel sheet No. ll, the C content and the N content were out of the range of the present invention, and the solid solution N amount and the martensite amount were out of the range of the present invention. Therefore, the BH amount, ΔTS decreased, and ΔΥ-E1 increased. are doing. In addition, in steel sheet No. 12, Al, NZA1, and N were out of the range of the present invention, the amount of solute N was out of the range of the present invention, and the average grain size of ferrite was out of the range of the present invention. Therefore, the BH amount and ΔTS decrease, and ΔΥ-El increases.
鋼板 No. 13 は、 スラブ加熱温度と F D Tが本発明の好適範囲を外れ、 固溶 N量、 マルテンサイ ト量が本発明範囲を低くはずれ、 フヱライ トの平均結晶粒径が本発明 の範囲を高く外れ、 そのため、 r値、 B H量、 A T Sが低下している。 また、 鋼板 N 0. 1 は、 熱延後の卷取り温度が本発明範囲を外れ、 固溶 N量が本発明範囲を低くは ずれ、 フェライ トの平均結晶粒径が本発明の範囲を高く外れ、 そのため、 r値、 B H量、 A T Sが低下している。  In the steel sheet No. 13, the slab heating temperature and the FDT were out of the preferred ranges of the present invention, the amount of solid solution N and the amount of martensite were out of the range of the present invention, and the average crystal grain size of the fiber was increased in the range of the present invention. The r value, BH amount, and ATS have decreased. Further, in the steel sheet N 0.1, the winding temperature after hot rolling was out of the range of the present invention, the amount of solute N departed from the range of the present invention, and the average crystal grain size of ferrite increased in the range of the present invention. The r value, BH amount, and ATS have decreased.
鋼板 No. 15 は、 連続焼鈍温度が本発明の好適範囲を外れ、 マルテンサイ トが生成 せず、 フェライ トの平均結晶粒径が本発明の範囲を高く外れ、 そのため、 B H量、 厶 T Sが低下し Δ Υ-Elが増加している。 また、 鋼板 No. 16 は、 箱焼鈍を実施せず、 望ましい集合組織が発達しないため、 特に r値が低下している。 またフェライ トの 平均粒径、 マルテンサイトの面積率も本癸明の範囲を外れている。  In the steel sheet No. 15, the continuous annealing temperature was out of the preferred range of the present invention, no martensite was formed, and the average grain size of ferrite was out of the range of the present invention, so that the BH content and the TS decreased. Δ Υ-El is increasing. In addition, for steel sheet No. 16, no box annealing was performed, and the desired texture was not developed, so the r-value was particularly low. The average particle size of ferrite and the area ratio of martensite are also out of the range of the present invention.
実施例 7 Example 7
表 1 8に示す組成になる鋼を、 実施例 1と同様の方法でスラブとなし、 該スラブ を表 1 9に示す条件で加熱し、 粗圧延して 30mm厚のシートパーとし、 ついで表 1 9 に示す条件の仕上圧延を施す熱間圧延工程により熱延板とした。 なお、 一部につい ては、 粗圧延後で仕上圧延入側で相前後するシートバー同士を溶融圧接法で接合し て連続圧延した。 また、 シートパーの幅端部、 長さ方向端部を誘導加熱方式のシー トパーエッジヒータ、 シートパーヒータを使用してシートパーの温度を調節した。 これら熱延板を酸洗および表 1 9に示す条件の冷間圧延からなる冷間圧延工程に より 1. 6 ram厚の冷延板と-した。 ついで、 これら冷延板に表 1 9に示す条件で、 箱焼 鈍、 ついで連続焼鈍炉による連続焼鈍を行った。 なお、 箱焼鈍の焼鈍温度はいずれ も再結晶温度以上とした。  A steel having the composition shown in Table 18 was formed into a slab in the same manner as in Example 1, and the slab was heated under the conditions shown in Table 19 and roughly rolled to form a 30 mm thick sheet par. A hot rolled sheet was obtained by a hot rolling step of performing finish rolling under the following conditions. In addition, a part of the sheet bars, which were adjacent to each other on the entry side of the finish rolling after the rough rolling, were joined by a melt pressure welding method and were continuously rolled. In addition, the width of the sheet par and the end in the length direction of the sheet par were controlled by using an induction heating type sheet par edge heater and a sheet par heater. These hot-rolled sheets were formed into 1.6 ram-thick cold-rolled sheets by a cold rolling process including pickling and cold rolling under the conditions shown in Table 19. Next, box annealing was performed on these cold-rolled sheets under the conditions shown in Table 19, and then continuous annealing was performed in a continuous annealing furnace. Note that the annealing temperatures of the box annealing were all higher than the recrystallization temperature.
得られた冷延焼鈍板について、 実施例 1と同様に固溶 N量、 微視組織、 引張特性、 r値測定、 歪時効硬化特性を調査した。 また、 各冷延焼鈍板の幅方向および長手方 向について、 各 10箇所で引張特性を調査し、 降伏強さ、 引張強さ、 伸びのばらつき を調査した。 なお、 ばらつきは測定した個所すベての中での最大値と最小値の差、 例えば S Y S = (Y Sの最大値） ― （Y Sの最小値） で表示した。 それらの結果 を表 2 0に示す。 With respect to the obtained cold-rolled annealed sheet, the amount of dissolved N, microstructure, tensile properties, r-value measurement, and strain age hardening properties were examined in the same manner as in Example 1. In addition, the tensile properties of each cold-rolled annealed sheet were examined at 10 locations in the width direction and the longitudinal direction, and the variation in yield strength, tensile strength, and elongation investigated. The variability was expressed as the difference between the maximum and minimum values of all measured points, for example, SYS = (maximum value of YS)-(minimum value of YS). Table 20 shows the results.
本発明例は、 いずれも優れた歪時効硬化特性と高い r値を有し、 製造条件の変動 にもかかわらず、 安定して格段に高い B H量、 A T S、 平均 r値を示した。 また、 本発明例では、 連続圧延とシートパーの長手方向、 幅方向温度調整を実施すること により、 製品鋼板の板厚精度および形状が向上し、 材質ばらつきが減少することを 確認した。 産業上の利用可能性  Each of the examples of the present invention had excellent strain age hardening characteristics and a high r-value, and showed a stably remarkably high BH content, ATS, and an average r-value despite fluctuations in production conditions. Further, in the present invention example, it was confirmed that by performing the continuous rolling and the temperature adjustment in the longitudinal direction and the width direction of the sheet par, the thickness accuracy and shape of the product steel sheet were improved, and the variation in the material was reduced. Industrial applicability
本発明によれば、 プレス成形時に優れた深絞り性を確保しつつ、 プレス成形-熱処 理により T Sが大きく増加する冷延鋼板が得られる。 この冷延鋼板から電気亜鉛め つき鋼板、 溶融亜鉛めつき鋼板、 合金化溶融亜鉛めつき鋼板を工業的に製造できる という優れた効果を奏する。 According to the present invention, it is possible to obtain a cold-rolled steel sheet in which TS is greatly increased by press forming-heat treatment while securing excellent deep drawability during press forming. From this cold-rolled steel sheet, there is an excellent effect that an electro-galvanized steel sheet, a hot-dip galvanized steel sheet, and an alloyed hot-dip galvanized steel sheet can be industrially manufactured.
銅 成 分 祖 成 （mass% ) Copper component (mass%)
記号 (1)'式 * (2)，式 ** 備 考 o C N Si Mn B Al Nb P S  Symbol (1) 'expression * (2), expression ** Remarks o C N Si Mn B Al Nb P S
A 0. o 0009 0. on 0.01 0.12 0.0009 0.010 0.016 0.009 0.005 0.0023 一 0.0001 適合例 A 0.o 0009 0.on 0.01 0.12 0.0009 0.010 0.016 0.009 0.005 0.0023 1 0.0001 Applicable example
B 0.0020 0.01 0.丄 1 0.012 0.035 0.006 0.0021 ― 0.0003B 0.0020 0.01 0. 丄 1 0.012 0.035 0.006 0.0021 ― 0.0003
C 0.0005 0.009 0.01 0， 09 0.0005 0.009 0.010 0.011 0.005 0.0022 C 0.0005 0.009 0.01 0, 09 0.0005 0.009 0.010 0.011 0.005 0.0022
o o  o o
D 0.0020 0.021 0.01 0.50 0.020 0.035 0.030 0.004 0.0035 一 0.0003 D 0.0020 0.021 0.01 0.50 0.020 0.035 0.030 0.004 0.0035 one 0.0003
E 0.0003 o U. OU U. \i 0.0006 0.099 0.045 0.0022 一 0.0003E 0.0003 o U.OU U. \ i 0.0006 0.099 0.045 0.0022 one 0.0003
F 0.0011 0.030 0.80 0.80 0.0011 0.028 0.025 0.009 0.005 0.0103 - 0.0005 "F 0.0011 0.030 0.80 0.80 0.0011 0.028 0.025 0.009 0.005 0.0103-0.0005 "
G 0.018 0.70 0.12 0.0008 0.012 0.008 0.005 0.0080 一 0.0001G 0.018 0.70 0.12 0.0008 0.012 0.008 0.005 0.0080 one 0.0001
H 0.0005 0.020 0.35 0.11 0.020 0.025 0.007 0.005 0.0041 ― 0.0011H 0.0005 0.020 0.35 0.11 0.020 0.025 0.007 0.005 0.0041 ― 0.0011
I 0.0098 0.002 0.01 0.12 0.0007 0.038 0.055 0.009 0.005 - 0.0269 0.0027 比較例I 0.0098 0.002 0.01 0.12 0.0007 0.038 0.055 0.009 0.005-0.0269 0.0027 Comparative example
J 0.0022 0.001 0.50 0.12 0.0008 0.012 0.001 0.008 0.005 一 0.0064 J 0.0022 0.001 0.50 0.12 0.0008 0.012 0.001 0.008 0.005 one 0.0064
CO  CO
K 0.0260 0.003 0.02 0.25 0.0001 0.035 0.001 0.013 0.007 -0.0154 0.0259 K 0.0260 0.003 0.02 0.25 0.0001 0.035 0.001 0.013 0.007 -0.0154 0.0259
L 0.0027 0. on 0.01 0.12 o 0.0007 0.014 0.001 0.009 0.005 0.0027 L 0.0027 0.on 0.01 0.12 o 0.0007 0.014 0.001 0.009 0.005 0.0027
o o 0.0026  o o 0.0026
o  o
* (1)，式 ： N %— (14/93 · Nb% + 1 o o4/27 ' Al% + 14/11 · B % ) (0.0015  * (1), equation: N% — (14/93 · Nb% + 1 o o4 / 27 'Al% + 14/11 · B%) (0.0015
o 以上が本発明の適合範囲） ** (2)，式 ： C % - (0.5 · 12/93 - Nb % ) ( 0以下が本発明の適合範西）  o The above is the applicable range of the present invention) ** (2), Formula: C%-(0.5 · 12/93-Nb%) (0 or less is the applicable range of the present invention)
o  o
o  o
o o o o
o  o
o  o
o o
o  o
o  o
o  o
o 表 2 o Table 2
Figure imgf000076_0001
Figure imgf000076_0001
* 仕上圧延終了温度が Ar3変態点未満。 仕上圧延後の冷却条件は、 冷却開始時間（ s )及び冷却速度（¾/s)を示す, * Finish rolling end temperature is lower than Ar 3 transformation point. The cooling conditions after finish rolling indicate the cooling start time (s) and the cooling rate (¾ / s),
表 3 Table 3
Figure imgf000077_0001
Figure imgf000077_0001
91 91
Figure imgf000078_0001
Figure imgf000078_0001
挲  挲
tOOlO/lOdT/lDd IC1706/10 OAV tOOlO / lOdT / lDd IC1706 / 10 OAV
6 ΐ 6 I L 6 ΐ ' 0 0 8 0 ' 06 ΐ 6 I L 6 ΐ '0 0 8 0' 0
S Ζ 0 Z I I z ■ 0 2 0 * 0S Ζ 0 Z I I z ■ 0 2 0 * 0
Ζ f ε g i ΐ ε · o 6 ^ 0 · 0Ζ f ε g i ΐ ε · o 6 ^ 0 · 0
9 S Z 9 L 2 ί- ' 0 9 ε 0 ' 09 S Z 9 L 2 ί- '0 9 ε 0' 0
8 9 S Z 9 Z · 0 0 2 0 * 08 9 S Z 9 Z0 0 2 0 * 0
Β d Κ B d H % S X V ¾ J X S 丄 ΐ V / N ΐ V Β d Κ B d H% S X V ¾ J X S 丄 ΐ V / N ΐ V
^OOlO/lOdf/X3d t 06/10 OAV 表 6 ^ OOlO / lOdf / X3d t 06/10 OAV Table 6
3 Three
QD
Figure imgf000080_0001
QD
Figure imgf000080_0001
添加されていない Nb, Ti, V, Bについては上の濃度を Γ- と表記してあり N/ (Al+Nb+Ti +V+B)の計算には濃度を 「0」 として取り扱う。 For non-added Nb, Ti, V, B, the above concentration is indicated as Γ-, and the concentration is treated as “0” in the calculation of N / (Al + Nb + Ti + V + B).
表 7 Table 7
to to
Figure imgf000081_0001
Figure imgf000081_0001
SRT:スラフ *カ微献、 RDT=赃延出側喊、 FBHthiEMA側鍵、 FmH±JbJ£延出側暖 CT=卷取鍵 SRT: Sluff * Fine offer, RDT = 赃 Extending sword, FBHthiEMA side key, FmH ± JbJ £ Extending side warm CT = Rewind key
1=冷麵反、 π= めっき ¾、 m=合金t«¾&めっき » 1 = Cold plating, π = Plating m, m = Alloy t «¾ & Plating»
表 8 固 溶 N 歪 時 効 処 理 前 の 引 張 特 性 歪 時 効 硬 化 特 性 Table 8 Tensile properties before solid solution N strain aging treatment Strain aging hardening properties
No. 鋼 Y S T S E 1 r値 T S x r値 B H A T S 備 考 in 1 a / M P a Pa l a  No. Steel Y S T S E 1 r value T S x r value B H A T S Remarks in 1 a / M P a Pa la
1 A 0 . 0 0 6 9 2 2 5 3 2 1 5 3 2 . 4 7 7 0 1 2 2 7 5 実 施 例 1 A 0 .0 0 6 9 2 2 5 3 2 1 5 3 2 .4 7 7 0 1 2 2 7 5
2 B 0 . 0 0 8 9 2 7 4 3 9 1 4 3 2 . 3 8 9 9 1 8 3 9 3 実 施 例2 B 0 .0 0 8 9 2 7 4 3 9 1 4 3 2 .3 8 9 9 1 8 3 9 3
3 C 0 . 0 0 5 4 2 2 1 3 1 6 5 4 2 . 8 8 8 5 9 7 7 2 実 施 例3 C 0 .0 0 5 4 2 2 1 3 1 6 5 4 2 .8 8 8 5 9 7 7 2
4 D 0 . 0 0 4 9 2 2 1 3 1 6 5 4 2 . 8 8 8 5 9 0 6 6 実 施 例4 D 0.04 9 2 2 1 3 1 6 5 4 2 .8 8 8 5 9 0 6 6
5 E 0 . 0 0 5 0 3 0 4 4 3 5 3 9 2 . 0 8 7 0 8 0 6 3 実 施 例5 E 0 .0 0 5 0 3 0 4 4 3 5 3 9 2 .0 8 7 0 8 0 6 3
6 F 0 . 0 0 8 8 3 0 4 4 3 4 3 9 2 . 1 9 1 1 1 3 3 8 7 実 施 例6 F 0 .0 0 8 8 3 4 4 3 4 3 9 2 .1 9 1 1 1 3 3 8 7
00 G 0 . 0 0 0 0 2 2 4 3 2 0 5 3 2 . 8 8 9 6 3 0 00 G 0 .0 0 0 0 2 2 4 3 2 0 5 3 2 .8 8 9 6 3 0
o 7 比 較 例 o 7 Comparative example
8 H 0 . 0 0 0 0 2 8 4 4 0 5 4 2 2 . 1 8 5 1 2 0 比 較 例 8 H 0 .0 0 0 0 0 2 8 4 4 0 5 4 2 2 .1 8 5 1 2 0 Comparative example
9 A 0 . 0 0 7 0 2 1 5 3 1 1 5 0 2 . 3 7 1 5 1 5 2 8 3 比 較 例9 A 0 .0 0 7 0 2 1 5 3 1 1 5 0 2 .3 7 1 5 1 5 2 8 3 Comparative example
1 0 B 0 . 0 0 8 2 2 7 4 3 9 1 4 3 1 . 9 7 4 3 1 4 3 9 4 比 較 例1 0 B 0 .0 0 8 2 2 7 4 3 9 1 4 3 1 .9 7 4 3 1 4 3 9 4 Comparative example
1 1 C 0 . 0 0 3 5 2 3 6 3 3 1 5 1 2 . 0 6 6 2 9 3 4 8 比 較 例1 1 C 0 .0 0 3 5 2 3 6 3 3 1 5 1 2 .0 6 6 2 9 3 4 8 Comparative example
1 2 D 0 . 0 0 4 1 2 4 ϊ 3 3 6 5 1 2 . 1 7* 0 6 • 8 5 4 9 比 較 例 1 2 D 0 .0 0 4 1 2 4 ϊ 3 3 6 5 1 2 .1 7 * 0 6 • 8 5 4 9 Comparative example
o o
2S00 ·0:Β3 2S00 ·0 OIOO ·0 0800 ,0 U ·ΐ 0210 ·0 iOO ·ο TOO ·0 800 ·0 18 ·0 81 ·0 9β00 ·0 ΐΟΟ ·0 6000 ·0 OiOO '0 -s ΟΟίΌ ·0 100 ·0 TOO ·0 ΑΟΟ Ό 28 ·0 LI Ό 9200 ·0 d 2S000: Β3 2S000 OIOO0 0800, 0 U0 · 0210 TO0 iOO ο TOO 00 800 00 18 00 81 00 9β00 00 ΐΟΟ0 60000 OiOO '0 -s s · 0 100 · 0 TOO · 0 ΑΟΟ Ό 28 · 0 LI Ό 9200 · 0 d
STOO ·0 δ^ΐθ '0 ΟΙΤ ·0 Ί 02Τ0 Ό Ζ,ΟΟ ·0 100 ·0 800 ·0 18 ·0 LI 'Ο S200 ·0 ΐΟΟΟ ·0 > 9000 ·0 0S00 ·0 ο ·ζ 02Τ0 '0 900 ·0 ZOO ·0 090 ·0 98 ·0 91 ·0 9Ζ00 ·0STOO0 δ ^ ΐθ '0 ΟΙΤ0 Ί 02Τ0 Ζ Ζ, ΟΟ 900 0 ZOO 0 090 0 98 98 0 91 0 00 0
ΤΟΟΟ Ό> 0000 ·0 ΐΟΟΟ ·0 ο ο ·0 S00 ·0 zoo ·0 090 Ό S8 ·0 SI ·0 S200 ·00000 Ό> 0000 · 0 ΐΟΟΟ · 0 ο ο · 0 S00 · 0 zoo · 0 090 Ό S8 · 0 SI · 0 S200 · 0
TCOO "0 ΖΤΟΟ '0 0600 Ό L0 ·0 0刚 *0 εοο ·ο 2Τ0 ·0 09 ·ΐ Ζ ·0 L0 ·0TCOO "0 ΖΤΟΟ '0 0600 Ό L0 00 0 刚 * 0 εοο ο 2Τ0 009 ΐ Ζ 00 L0
30 ·0:ΐΝ 'ΟΙ '0："3 0200 ·0 ΙΙΟΟ ·0 9 00 ·0 Of ·ζ ΟΖΐΟ ·0 900 '0 ZOO Ό 800 ·0 ΙΖ ·χ 51 ·0 ZZQ ·030 · 0: ΐΝ 'ΟΙ' 0: "3 0200 · 0 ΙΙΟΟ · 0 9 00 · 0 Of · ζ ΟΖΐΟ · 0 900 '0 ZOO Ό 800 · 0 ΙΖ · χ 51 · 0 ZZQ · 0
ΖΟ "0:Λ '30'0:ΐΝ 8100 ·0 ΟΐΟΟ ·0 6800 ·0 9Ε OSTO ·0 πο ·0 εοο ·ο goo 'ο 5Ζ Ί ST ·0 ΨΖΟΟ ·0 ΖΟ "0: Λ '30'0: ΐΝ 8100 · 0 ΟΐΟΟ · 0 6800 · 0 9Ε OSTO · 0 πο · 0 εοο · ο goo' ο 5Ζ Ί ST · 0 ΨΖΟΟ · 0
Ψ 0 '0 8000 *0 0900 ·0 LZ ·\ Ο^ΐΟ ·0 ΠΟ ·0 TOO ,0 800 ·0 Ζ 'ΐ 92 ·0 9200 ·0 ειο ·ο:ίΐ 9100 ·0 6000 ·0 ( 00 ·0 9Ζ, 'Ζ οπο '0 00 '0 TOO ·0 OO Ό 92 'I 91 ·0 S200 ·0 Ψ 0 '0 8000 * 0 0900 · 0 LZ · \ Ο ^ ΐΟ · 0 ΠΟ · 0 TOO, 0 800 · 0 Ζ' ΐ 92 · 0 9200 · 0 ειο · ο: ίΐ 9100 · 0 6000 · 0 (00 0 9Ζ, 'Ζ οπο' 0 00 '0 TOO 0 OO Ό 92' I 91 0 S2000
51 ·0:。Η 0200 ·0 0100 ·0 0800 ·0 Ί ΟΖΙΟ Ό ZOO ·0 TOO ·0 9^0 ·0 90 ·0 εζοο ·ο 51 · 0 :. Η 0200 · 0 0100 · 0 0800 · 0 Ί ΟΖΙΟ Ό ZOO · 0 TOO · 0 9 ^ 0 · 0 90 · 0 εζοο
ΟΐΟΟ ·0 6000 ·0 οζοο ·0 60 Ί 0210 ·0 ΠΟ '0 200 Ό 500 ·0 10 Ό ΖΖΟΟ ·0 ΟΐΟΟ · 0 6000 · 0 οζοο · 0 60 Ί 0210 · 0 ΠΟ '0 200 Ό 500 · 0 10 Ό ΖΖΟΟ · 0
2100 ·0 ετοο ·ο 2010 ·0 00 ·ΐ ΟΟΐΟ ·0 οτο ·0 300 ·0 010 ·0 88 ·0 91 ·0 OSTO ·02100 · 0 ετοο · ο 2010 · 0 00 · ΐ ΟΟΐΟ · 0 οτο · 0 300 · 0 010 · 0 88 · 0 91 · 0 OSTO
STOO '0 9000 ·0 0300 '0 οο ·ε 0210 '0 刚 Ό TOO ·0 0 0 ·0 9ε Ό ζοο ·0STOO '0 9000 · 0 0300' 0 οο · ε 0210 '0 刚 Ό TOO · 0 0 0 · 0 9ε Ό ζοο · 0
6000 ·0 6000 ·0 OiOO ·0 88 Ί 0910 ·0 800 Ό TOO ·0 00 '0 OS 19 ·0 0S00 ·06000 00 6000 00 OiOO 00 88 Ί 0910 00 800 Ό TOO 00 00 '0 OS 19 00 S00 00
SIOO "0 SOOO '0 0刚 ·0 Of 'τ 0310 *0 S00 ·0 zoo ·0 0S0 ·0 38*0 SI ·0 S200 ·0 ¾■ a qN Ζ6/ΖΙ qN TV/N N TV SIOO "0 SOOO '0 0 刚 0 Of' τ 0310 * 0 S00
(%薔葛) ^ ^ ^ (% Rose) ^ ^ ^
6 6
表 1 0 Table 10
00 00
Figure imgf000084_0001
Figure imgf000084_0001
Figure imgf000085_0001
Figure imgf000085_0001
AF ァシキユラ' -フ ライ 卜 溶融亜鉛め- き処理あり AF Ashkyura '-Flite Hot-dip galvanized
さ/ 6 O：AV· / 6 O: AV ·
o o
Figure imgf000086_0001
Figure imgf000086_0001
z x zx
表 1 3 Table 13
Figure imgf000087_0001
Figure imgf000087_0001
*) 潤滑圧延実施  *) Lubrication rolling
00 *) ·潤'滑圧延実施， シートバーヒータ、 エッジヒーター使用 00 *) · Smooth rolling is performed, using sheet bar heater and edge heater
表 14 鋼板 鋼 鋼板 鋼板組織 製品板特性 製品板板特性 歪時効硬化特性 耐時効性 備 考Table 14 Steel plate Steel plate Steel plate structure Product plate characteristics Product plate characteristics Strain aging hardening characteristics Aging resistance Remarks
No. . No. 固溶 N No.. No. Solid solution N
量 フ ライ卜 第 2相 引張特性 引張特性のばらつき * BH量 ATS  Amount Light Phase 2 Tensile properties Variation in tensile properties * BH amount ATS
囬¾率 粒径 AF YS TS E 1 平均 δ YS δ TS δ E 1  Rate Particle size AF YS TS E 1 Average δ YS δ TS δ E 1
面積率 r値  Area ratio r value
質量％ % μ m % MPa MPa % MPa MPa % MPa MPa %  Mass%% μ m% MPa MPa% MPa MPa% MPa MPa%
2-1 N 0.0120 91 7 9 315 445 37 1.6 10 10 5 15 15 0 本発明例 2-1 N 0.0120 91 7 9 315 445 37 1.6 10 10 5 15 15 0 Example of the present invention
2-2 N 0.0120 91 7 9 319 447 37 1.6 5 5 2 7 6 0 本発明例2-2 N 0.0120 91 7 9 319 447 37 1.6 5 5 2 7 6 0 Example of the present invention
2-3 N 0.0125 92 7 8 318 450 38 1.5 3 5 1 8 5 0 本発明例 2-3 N 0.0125 92 7 8 318 450 38 1.5 3 5 1 8 5 0 Example of the present invention
00 ) SYS, STS, δΕΙ: = (最大値—最小値） 00) SYS, STS, δΕΙ: = (maximum value-minimum value)
) AF：ァシキユラ一フェライト、 M：マルテンサイト、 B：べィナイト、 P：パーライト  ) AF: Ashura ferrite, M: Martensite, B: Bainite, P: Pearlite
<
Figure imgf000089_0002
<
Figure imgf000089_0002
Figure imgf000089_0001
Figure imgf000089_0001
s τ ¾ s τ ¾
表 1 6 Table 16
00 00
00 00
Figure imgf000090_0001
Figure imgf000090_0001
*) 潤滑圧延実施  *) Lubrication rolling
**) 冷却停止温度以下 350 ¾以上の滞留時間 ***) 温片装入  **) Dwell time above cooling stop temperature and above 350¾ ***)
****) 溶融亜鉛めつきの後に伸び率 0. 5 %の調質圧延 ****) Temper rolling with 0.5% elongation after hot-dip galvanizing
表 1 7 Table 17
0000
O O
Figure imgf000091_0001
Figure imgf000091_0001
Μ：マルテンサイ ト、 Β ：ペイナイ ト、 Ρ ：パ一ライ ト Μ: Martensite, Β: Paynight, Ρ: Private
表 1 8 鋼 化 学 成 分 （質量％) 変態点 （で）Table 18 Steel chemical composition (% by mass) Transformation point (D)
No. No.
C . Si Mn P S Al N N /Al その他 A ci A c3 C. Si Mn PS Al NN / Al Other A ci A c 3
M 0.052 0.01 0.60 0.035 0.001 0.002 0.0125 735 855 M 0.052 0.01 0.60 0.035 0.001 0.002 0.0125 735 855
CD CD
t t
表 1 9 Table 19
COCO
Figure imgf000093_0001
Figure imgf000093_0001
*) 潤滑圧延実施  *) Lubrication rolling
*) シートパ一ヒーター、 エッジヒータ一使用 *) Uses sheet heater and edge heater
表 2 0 Table 20
Figure imgf000094_0001
Figure imgf000094_0001
) M： マルテンサイ ト、 B ：べィナイ ト、 P ：パーライ ト ) M: Martensite, B: Venite, P: Pearlite
CO CO
t t

Claims

請 求 の 範 囲 The scope of the claims
1 · mass%で、  1 mass%,
C ： 0.15%以下、  C: 0.15% or less,
Si ： 1.0%以下、  Si: 1.0% or less,
Mn ： 2.0%以下、  Mn: 2.0% or less,
P ： 0.1%以下、  P: 0.1% or less,
S ： 0.02%以下、  S: 0.02% or less,
A1 ： 0.005 〜0.030%、  A1: 0.005 to 0.030%,
N ： 0.0050 〜0.0400%、  N: 0.0050-0.0400%,
を含み、 かつ N/A1： 0.30以上、 And N / A1: 0.30 or more,
固溶状態の Nが 0.0010%以上あり、 0.0010% or more of N in solid solution state,
残部が Feおよび不可避的不純物からなる組成を有することを特徴とする、 歪時効硬 化特性に優れた冷延鋼板。 A cold-rolled steel sheet excellent in strain aging hardening characteristics, characterized in that the balance has a composition consisting of Fe and unavoidable impurities.
2. 前記組成に加えてさらに、 mass%で、 下記 a群〜 d群のうちの 1群または 2 群以上を含むことを特徴とする、請求項 1記載の歪時効硬化特性に優れた冷延鋼板。 2. In addition to the composition, further comprising, in mass%, one or more of the following groups a to d, wherein the cold rolling is excellent in strain age hardening characteristics according to claim 1. steel sheet.
記  Record
a群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1.0%以下、  Group a: One or more of Cu, Ni, Cr and Mo are 1.0% or less in total,
b群： Nb、 Ti、 Vのうちの 1種または 2種以上を合計で 0.1%以下  Group b: 0.1% or less of one or more of Nb, Ti and V
c群： Bを 0.0030%以下  Group c: B is 0.0030% or less
d群： Ca、 REM の 1種または 2種を合計で 0.0010〜0.010%  Group d: Ca or REM 1 or 2 types in total 0.0010 to 0.010%
3. nmss%にて 3.At nmss%
C ： 0.01%未満、 C: less than 0.01%,
Si： 0.005 〜1.0%、 Si: 0.005 to 1.0%,
Mn： 0.01〜： I.5%、 P ： 0.1%以下、 Mn: 0.01 ~: I.5%, P: 0.1% or less,
S ： 0.01%以下、  S: 0.01% or less,
Al： 0.005 -0.030%,  Al: 0.005 -0.030%,
N： 0.005 〜0.040%,  N: 0.005 to 0.040%,
を含み、 かつ N/A1： 0.30以上、 And N / A1: 0.30 or more,
固溶状態の Nが 0.0010%以上あり、 0.0010% or more of N in solid solution state,
残部が Feおよび不可避的不純物からなる組成を有することを特徴とする歪時効硬化 特性に優れた冷延鋼板。 A cold-rolled steel sheet having excellent strain aging hardening characteristics, characterized in that the balance is composed of Fe and unavoidable impurities.
4. 上記組成に加えてさらに、 mass%で 4. In addition to the above composition,
B ： 0.0001〜0.0030%、 B: 0.0001-0.0030%,
Nb： 0.005 -0.050%, Nb: 0.005 -0.050%,
を、 次式（1)， （2) With the following equations (1) and (2)
N%≥0.0015 + 14/93 - Nb% + 14/27 - Al% + 14/11 · B% —- (1)  N% ≥0.0015 + 14/93-Nb% + 14/27-Al% + 14 / 11B% --- (1)
C%≤ 0.5 - (12/93) - Nb% (2) を満足する範囲において含有し、残部は実質的に Feの組成になることを特徴とする、 請求項 3に記載の歪時効硬化特性に優れた冷延鋼板。  The strain-age hardening characteristic according to claim 3, characterized in that it is contained within a range satisfying C% ≤ 0.5-(12/93)-Nb% (2), and the balance is substantially Fe. Excellent cold rolled steel sheet.
5. 上記組成に加えてさらに、 mass%で必要に応じて Cu、 Ni、 Moのうちの 1種ま たは 2種以上を合計で 1.0 %以下含むことを特徴とする請求項 3または 4に記載の 歪時効硬化特性に優れた冷延鋼板。 5. The method according to claim 3, wherein in addition to the above composition, one or more of Cu, Ni, and Mo are further contained in a mass% of 1.0% or less as necessary. A cold-rolled steel sheet having excellent strain age hardening characteristics as described.
6. 鋼板の結晶粒径が 20 /xm 以下であることを特徴とする請求項 1から 5に記 載の歪時効硬化特性に優れた冷延鋼板。 6. The cold-rolled steel sheet having excellent strain aging hardening characteristics according to claim 1, wherein the crystal grain size of the steel sheet is 20 / xm or less.
7. 熱処理温度： 120 -200 ¾の低温域にて、 成形後の強度上昇代： 60 MPa以上 を有する請求項 1から 6に記載の歪時効硬化特性に優れた冷延鋼板。 7. Heat treatment temperature: In the low temperature range of 120-200¾, strength increase after molding: 60 MPa or more 7. The cold-rolled steel sheet according to claim 1, which has excellent strain aging hardening characteristics.
8. 請求項 1から 7に記載の冷延鋼板の表面に、 電気亜鉛めつき、 溶融亜鉛めつ き、 および合金化溶融亜鉛めつき層をそなえることを特徴とする、 歪時効硬化特性 に優れた電気亜鉛めつき、 溶融亜鉛めつき、 および合金化溶融亜鉛めつき鋼板。 8. The surface of the cold-rolled steel sheet according to claim 1 is provided with an electro-zinc plating, a hot-dip zinc plating, and an alloyed hot-dip galvanized layer, and has excellent strain aging hardening characteristics. Steel galvanized, hot-dip galvanized, and alloyed hot-dip galvanized steel sheets.
9. mass%にて 9. At mass%
C ： 0.01%未満、 C: less than 0.01%,
Si： 0.005 〜1.0%、  Si: 0.005 to 1.0%,
Mn： 0.01〜1.5%、  Mn: 0.01-1.5%,
P ： 0.1%以下、  P: 0.1% or less,
S ： 0.01%以下、  S: 0.01% or less,
A1： 0.005 〜0.030%、  A1: 0.005 to 0.030%,
N： 0.005 〜0.040%、  N: 0.005 to 0.040%,
を含み、 かつ N/A1： 0.30以上 And N / A1: 0.30 or more
を満足する範囲において含有し、 残部は実質的に Feの組成になる鋼片を、 熱間圧延 し、 その際、 仕上圧延終了後直ちに冷却を開始して卷取り温度： 400 〜800 で卷 取り、 その後圧下率： 60〜95%の冷間圧延を施したのち、 650~900での温度で再結 晶焼鈍を施すことを特徴とする、 歪時効硬化特性に優れた冷延鋼板の製造方法。 The steel slab having a composition of substantially Fe is hot-rolled. At that time, cooling is started immediately after finishing rolling, and the coil is wound at a winding temperature of 400 to 800. A method for producing a cold-rolled steel sheet having excellent strain aging hardening characteristics, wherein cold rolling is performed at a reduction ratio of 60 to 95%, and then recrystallization annealing is performed at a temperature of 650 to 900. .
10. 上記組成に加えてさらに、 mass%で 10. In addition to the above composition,
B ： 0.0001〜0.0030%、 B: 0.0001-0.0030%,
Nb： 0.005 〜0.050%、  Nb: 0.005 to 0.050%,
を、 次式（1)， (2) With the following equations (1) and (2)
N%≥0.0015 + 14/93 · Nb% + 14/27 · Al% + 14/11 · B% --- (1) C%≤ 0.5 · (12/93) - Nb% --- (2) を満足する範囲において含有し、残部は実質的に Feの組成になることを特徴とする、 請求項 9に記載の歪時効硬化特性に優れた冷延鋼板の製造方法。 N% ≥ 0.0015 + 14 / 93Nb% + 14 / 27Al% + 14 / 11B% --- (1) C% ≤ 0.5 (12/93)-Nb% --- (2) 10. The method for producing a cold-rolled steel sheet having excellent strain aging hardening characteristics according to claim 9, wherein the content is in a range satisfying the following, and the balance is substantially a Fe composition.
1 1. 上記した再結晶焼鈍における昇温過程において、 500 から再結晶温度ま での温度域を 1〜20で/ sの速度で昇温することを特徴とする、 請求項 9または 10 に記載の歪時効硬化特性に優れた冷延鋼板の製造方法。 11. The temperature rise process in the recrystallization annealing described above, wherein a temperature range from 500 to a recrystallization temperature is raised at a rate of 1 to 20 / s at a rate of 1 to 20. For producing cold rolled steel sheets with excellent strain age hardening characteristics.
12. 請求項 9から 1 1において、 再結晶焼鈍後、 溶融亜鉛めつき処理、 ついで 加熱合金化処理を施すこと特徴とする、 歪時効硬化特性に優れた合金溶融亜鉛めつ き鋼板の製造方法。 12. The method for producing a galvanized steel sheet having excellent strain aging hardening characteristics according to claims 9 to 11, wherein after recrystallization annealing, a molten zinc plating treatment and then a heat alloying treatment are performed. .
13. mass%で、 13. mass%,
C ： 0.01%未満、 C: less than 0.01%,
Si： 0.005 —1.0%,  Si: 0.005 -1.0%,
Mn： 0.01〜： 1.5%、  Mn: 0.01 ~: 1.5%,
P ： 0.1%以下、  P: 0.1% or less,
S ： 0.01%以下、  S: 0.01% or less,
A1： 0.005 〜0.030%、  A1: 0.005 to 0.030%,
N： 0.005 —0.040%、  N: 0.005 -0.040%,
を含み、 かつ N/A1： 0.30以上、 And N / A1: 0.30 or more,
固溶状態の Nが 0.0010%以上あり、 0.0010% or more of N in solid solution state,
残部が Feおよび不可避的不純物からなる組成を有し、 The balance has a composition consisting of Fe and unavoidable impurities,
TS X r値： 750MPa以上であることを特徴とする、 歪時効硬化特性に優れた深絞り 用冷延鋼板。  TS X r value: Cold rolled steel sheet for deep drawing with excellent strain age hardening characteristics, characterized by being 750 MPa or more.
14. 上記組成に加えてさらに、 mass%で B ： 0.0001〜0.0030%、 14. In addition to the above composition, B: 0.0001-0.0030%,
Nb： 0.005 〜0·050%、 Nb: 0.005 to 0.050%,
を、 次式（1)， （2) With the following equations (1) and (2)
Ν%≥0.0015 + 14/93 - Nb% + 14/27 - Al% + 14/11 - B% -― (1) C%≤ 0.5 · (12/93) - Nb% ―— (2) を満足する範囲において含有し、  Ν% ≥0.0015 + 14/93-Nb% + 14/27-Al% + 14/11-B% ---- (1) Satisfies C% ≤ 0.5 · (12/93)-Nb% ---- (2) To the extent that
残部が Feおよび不可避的不純物からなる組成を有し、 The balance has a composition consisting of Fe and unavoidable impurities,
TSX r値： 750MPa以上であることを特徴とする、 請求項 1 3に記載の歪時効硬化 特性に優れた深絞り用冷延銅板。  The cold-rolled copper sheet for deep drawing having excellent strain aging hardening characteristics according to claim 13, characterized in that the TSX r value is 750 MPa or more.
15. 請求項 13記載の鋼組成に加えてさらに、 mass %で 15. In addition to the steel composition according to claim 13,
B ： 0.0001〜0.0030%、 B: 0.0001-0.0030%,
Nb： 0.005 〜0.050%、 Nb: 0.005 to 0.050%,
Ti： 0.005〜0.070%、 Ti: 0.005-0.070%,
V： 0.005 〜0.10%のうち 1種または 2種以上を含有し、  V: One or more of 0.005 to 0.10%
かつ N/(A1+Nb+Ti+V+B) ： 0.30以上、 And N / (A1 + Nb + Ti + V + B): 0.30 or more,
固溶状態の Nが 0.0010%以上あり、 0.0010% or more of N in solid solution state,
残部 Feおよび不可避的不純物からなる組成を有し、 The balance has a composition consisting of Fe and unavoidable impurities,
TS X r値： 750MPa以上であることを特徴とする、 請求項 1 3に記載の歪時効硬化 特性に優れた深絞り用冷延鋼板。  14. The cold-rolled steel sheet for deep drawing according to claim 13, which has a TS Xr value of 750 MPa or more.
10. mass%で、 10. mass%,
C ： 0.01%未満、 C: less than 0.01%,
Si： 0.005 -1.0%,Si: 0.005 -1.0%,
n： 0.01〜1.5%、  n: 0.01-1.5%,
P ： 0.1%以下、 S ： 0.01%以下、 P: 0.1% or less, S: 0.01% or less,
Al： 0.005 〜0.030%、 Al: 0.005 to 0.030%,
N： 0,005 -0.040%, N: 0,005 -0.040%,
を含み、 Including
B ： 0.0001-0.0030%、  B: 0.0001-0.0030%,
Nb： 0.005 〜0·050%、 Nb: 0.005 to 0.050%,
Ti： 0.005 〜0.070%、 Ti: 0.005 to 0.070%,
V： 0.005 〜0.10%のうち 1種または 2種以上を含み、 かつ  V: contains one or more of 0.005 to 0.10%, and
N/(A1+Nb+Ti+V+B) ： 0,30以上になる組成を有する鋼素材を、  N / (A1 + Nb + Ti + V + B): A steel material having a composition of 0.30 or more,
950 以上に加熱後、 粗圧延終了温度を 1000で以下 Ar3以上として粗圧延し 続いて Ar 3以下 600 で以上の温度域で潤滑しつつ仕上圧延して卷き取り、 その際粗圧延開始から仕上圧延終了までの全圧下率を 80%以上とし、  After heating to 950 or more, rough rolling is performed at a rough rolling end temperature of 1000 or less and Ar3 or more, followed by finish rolling while lubricating at a temperature of Ar3 or less and 600 or more in the above temperature range. The total rolling reduction until the end of rolling is 80% or more,
得られた熱延板を再結晶焼鈍し、 Recrystallization annealing of the obtained hot rolled sheet,
次いで圧下率 60~95%で冷間圧延し、 Then cold-rolled at a rolling reduction of 60 to 95%,
得られた冷延板を再結晶焼鈍することを特徴とする歪時効硬化特性に優れた深絞り 用冷延鋼板の製造方法。 A method for producing a cold-rolled steel sheet for deep drawing having excellent strain aging hardening characteristics, characterized by recrystallization annealing the obtained cold-rolled sheet.
1 7. mass%で、 1 7. mass%,
C ： 0.0015〜0.025%、  C: 0.0015-0.025%,
Si： 1.0 %以下、  Si: 1.0% or less,
Mn： 2.0 %以下、  Mn: 2.0% or less,
P ： 0.1%以下、  P: 0.1% or less,
S ： 0.02%以下、  S: 0.02% or less,
A1 : 0.02.%以下、  A1: 0.02% or less,
N： 0.0050〜0.0250%、  N: 0.0050-0.0250%,
を含み、 つ B： 0. 0001〜0. 0050%、 Including B: 0.0001 to 0.0050%,
Nb： 0. 002 ~0. 050%, Nb: 0.002 ~ 0.050%,
の 1種または 2種以上を含み、 かつ N/A1が 0. 3 以上、 固溶状態としての Nを 0. 001 0%以上含有し、 残部が Feおよび不可避的不純物からなる組成と、 A composition comprising at least 0.3 of N / A1 and at least 0.001% of N as a solid solution, with the balance being Fe and unavoidable impurities,
面積率で 5 %以上のァシキユラ一フェライ ト相と平均結晶粒径： 20 / m 以下のフエ ライト相から成る組織を有し、 It has a structure consisting of a ferrite phase with an area ratio of 5% or more and a ferrite phase with an average crystal grain size of 20 / m or less.
r値： 1. 2 以上であることを特徴とする成形性、 歪時効硬化特性おょぴ耐常温時効 性に優れた冷延鋼板。  r-value: A cold-rolled steel sheet with excellent formability and strain age hardening characteristics, characterized by being 1.2 or more, and excellent room temperature aging resistance.
1 8 . 前記組成に加えてさらに、 raass%で、 下記 a群〜 c群のうちの 1群または 2群以上を含むことを特徴とする請求項 1 7に記載の冷延鋼板。 18. The cold-rolled steel sheet according to claim 17, further comprising one or more of the following groups a to c in raass% in addition to the composition.
記  Record
a群： Cu、 、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0 %以下 b群： Ti、 Vのうちの 1種または 2種を合計で 0. 1 %以下  Group a: 1.0% or less in total of one or more of Cu, Cr, and Mo Group b: 0.1% or less in total of one or two of Ti and V
c群： Ca、 REM の 1種または 2種を合計で 0. 0010~0. 010 %  Group c: One or two of Ca and REM in total 0.0010 to 0.010%
1 9 . mass%で、 1 9 mass%,
C： 0. 0015〜0. 025%、  C: 0.0015 to 0.025%,
Si： 1. 0%以下、  Si: 1.0% or less,
Mn： 2. 0%以下、  Mn: 2.0% or less,
P ： 0. 1%以下、  P: 0.1% or less,
S ： 0. 02%以下、  S: 0.02% or less,
A1： 0. 02%以下、  A1: 0.02% or less,
N： 0. 0050〜 0. 0250%、  N: 0.0050 to 0.0250%,
を含み、 かつ Including, and
B： 0. 0001〜0. 0050%、 Nb： 0. 002 〜0. 050%、 B: 0.0001 to 0.0050%, Nb: 0.002 to 0.050%,
の 1種または 2種以上を含み、 かつ N ZA1が 0. 3 以上である組成の鋼スラブを、 スラブ加熱温度： 1000°C以上に加熱し、 A steel slab containing one or more of the following and having a NZA1 of 0.3 or more is heated to a slab heating temperature of 1000 ° C or more,
粗圧延してシートバ^"となし、 Rough rolling and sheet bar ^ "
該シートパーに仕上圧延出側温度 '· 800 °C以上とする仕上圧延を施し、 The sheet par is subjected to finish rolling at a finish rolling exit temperature of 800 ° C or higher,
巻取温度： 800 以下で卷き取り熱延板とする熱間圧延工程と、 A hot rolling step of forming a hot rolled sheet at a winding temperature of 800 or less;
該熱延板に酸洗および冷間圧延を施し冷延板とする冷間圧延工程と、 A cold rolling step in which the hot-rolled sheet is subjected to pickling and cold rolling to form a cold-rolled sheet,
該冷延板にフユライトーオーステナイ トニ相域内の温度で連続焼鈍を行い、Continuous annealing was performed on the cold-rolled sheet at a temperature in the range of the Fulite-Austenitic Toni phase,
500 °C以下の温度域まで冷却速度： 10〜 300°C/sで冷却する冷延板焼鈍工程とを、 順次施すことを特徴とする Cooling speed to a temperature range of 500 ° C or less: Cold-rolled sheet annealing process of cooling at 10-300 ° C / s
r値： 1. 2 以上を有し、 成形性、 歪時効硬化特性おょぴ耐常温時効性に優れた冷延 鋼板の製造方法。  r-value: A method for producing cold-rolled steel sheets with a moldability of more than 1.2, and excellent strain-age hardening properties and room-temperature aging resistance.
2 0 . 前記組成に加えてさらに、 mass%で、 下記 a群〜 c群のうちの 1群または 2群以上を含むことを特徴とする請求項 1 9に記載の冷延鋼板の製造方法。 20. The method for producing a cold-rolled steel sheet according to claim 19, further comprising one or more of the following groups a to c in mass% in addition to the composition.
記 .  Record .
a群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0%以下、 b群： Ti、 Vのうちの 1種または 2種以上を合計で 0. 1%以下  Group a: 1.0% or less in total of one or more of Cu, Ni, Cr, and Mo; Group b: 0.1% in total of one or more of Ti, V Less than
c群： Ca、 REM の 1種または 2種を合計で 0. 0010〜0. 010%  Group c: Ca or REM 1 or 2 types in total 0.0010 to 0.010%
2 1 . mass%で、 2 1 mass%,
C ： 0. 025 〜0. 15% C: 0.025 to 0.15%
Si： 1. 0%以下、 Si: 1.0% or less,
Mn： 2. 0%以下、 Mn: 2.0% or less,
P ： 0. 08%以下、 P: 0.08% or less,
S ： 0. 02%以下、 Al： 0.02%以下、 S: 0.02% or less, Al: 0.02% or less,
N ： 0.0050-0.0250%  N: 0.0050-0.0250%
を含み、 かつ N/A1が 0.3 以上、 固溶状態としての Nを 0.0010%以上含有し、 残部 が Feおよぴ不可避的不純物からなる組成と、 And N / A1 is 0.3 or more, N is contained as a solid solution in 0.0010% or more, and the balance is Fe and unavoidable impurities.
平均結晶粒径： 10/zm 以下のフェライ ト相を面積 *で 80%以上含み、 さらに第 2相 として面積率で 2%以上のマルテンサイ ト相を含む組織とを有し、 Average grain size: has a structure containing at least 80% by area * of a ferrite phase of 10 / zm or less and a structure containing a martensite phase by area of 2% or more as a second phase,
r値： 1.2 以上であることを特徴とする高 r値と優れた歪時効硬化特性および常温 非時効性を有する高張力冷延鋼板。  r-value: A high-tensile cold-rolled steel sheet having a high r-value of 1.2 or more, excellent strain-age hardening characteristics and non-aging at room temperature.
22. 前記組成に加えてさらに、 mass%で、 下記 d群〜 g群のうちの 1群または 2群以上を含むことを特徴とする請求項 21に記載の高張力冷延鋼板。 22. The high-tensile cold-rolled steel sheet according to claim 21, further comprising, in addition to the composition, one or more of the following groups d to g in mass%.
記  Record
d群： Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1.0%以下 e群： Nb、 Ti、 Vのうちの 1種または 2種以上を合計で 0.1%以下  Group d: 1.0% or less in total of one or more of Cu, Ni, Cr, and Mo Group e: 0.1% or less in total of one or more of Nb, Ti, V
f 群： Bを 0.0030%以下  Group f: 0.0030% or less for B
g群： Ca、 REM の 1種または 2種を合計で 0.0010〜0.010%  g group: Ca or REM 1 or 2 types in total 0.0010 to 0.010%
23. mass%で、 23. mass%,
C ： 0.025 〜0.15% C: 0.025 to 0.15%
Si： 1.0%以下、 Si: 1.0% or less,
Mn： 2.0%以下、 Mn: 2.0% or less,
P ： 0.08%以下、 P: 0.08% or less,
S ： 0.02%以下、 S: 0.02% or less,
A1 : 0.02%以下、 A1: 0.02% or less,
N ： 0.0050〜0.0250% N: 0.0050-0.0250%
を含み、 かつ NZA1が 0.3 以上である組成の鋼スラブを、 スラブ加熱温度： 1000°C以上に加熱し、 And a steel slab having a composition of which NZA1 is 0.3 or more, Slab heating temperature: Heat to more than 1000 ° C,
粗圧延してシートパーとなし、 Rough rolling and sheet par, none
該シートパーに仕上圧延出側温度： 800 以上とする仕上圧延を施し、 Subjecting the sheet par to finish rolling at a finish rolling exit side temperature of 800 or more;
卷取温度： 800 t以下で卷き取り熱延板とする熱間圧延工程と、 Winding temperature: hot rolling process to make hot rolled sheet at 800 t or less,
該熱延板に酸洗およぴ冷間圧延を施し冷延板とする冷間圧延工程と、 A cold rolling step of subjecting the hot rolled sheet to pickling and cold rolling to form a cold rolled sheet;
該冷延板に焼鈍温度：再結晶温度以上 800 以下で箱焼鈍を施し、 Annealing the cold-rolled sheet at an annealing temperature of not less than the recrystallization temperature and not more than 800,
ついで焼鈍温度 ·· Ac 1変態点〜 （A c 3変態点— 20 ） で連続焼鈍を行い、 その後 500 以下の温度域まで冷却速度： 10~ 300で/5で冷却する冷延板焼鈍工程 とを、 順次施すことを特徴とする r値： 1. 2 以上の高 r値と優れた歪時効硬化特性 および常温非時効性を有する高張力冷延鋼板の製造方法。 Then annealing temperature · · Ac 1 transformation point - (A c 3 transformation point - 20) performs a continuous annealing at a cooling rate until further 500 following temperature range: 10 to the cold rolled sheet annealing step of cooling at 300/5 A method for producing a high-tensile cold-rolled steel sheet having a high r-value of 1.2 or more, excellent strain aging hardening characteristics and non-aging at room temperature.
2 4 . 前記連続焼鈍後の冷却に引き続いて、 前記冷却の冷却停止温度以下 350 以上の温度域で滞留時間 20 s以上の過時効処理を行うことを特徴とする請求項 2 3 に記載の高張力冷延鋼板の製造方法。 24. Following the cooling after the continuous annealing, an overaging treatment of a residence time of 20 s or more is performed in a temperature range of 350 or less of the cooling stop temperature of the cooling or less, wherein the high aging treatment according to claim 23, Manufacturing method of tension cold rolled steel sheet.
2 5 . 前記組成に加えてさらに、 mass %で、 下記 d群〜 g群のうちの 1群または 2群以上を含むことを特徴とする請求項 2 3または 2 4に記載の高張力冷延鋼板の 製造方法。 25. The high-tensile cold-rolled steel according to claim 23, further comprising, in addition to the composition, in mass%, one or more of the following groups d to g. Steel plate manufacturing method.
記  Record
d群: Cu、 Ni、 Cr、 Moのうちの 1種または 2種以上を合計で 1. 0%以下、  Group d: Cu, Ni, Cr, Mo or more of one or more of them, 1.0% or less in total,
e群： Nb、 Ti、 Vのうちの 1種または 2種以上を合計で 0. 1 %以下  Group e: 0.1% or less in total of one or more of Nb, Ti, and V
f 群： Bを 0. 0030%以下  f group: B is less than 0.0030%
g群： Ca、 RE の 1種または 2種を合計で 0. 0010〜0. 010%  g group: Ca or RE 1 or 2 types in total 0.0010 to 0.010%
PCT/JP2001/001004 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same WO2001090431A1 (en)

Priority Applications (6)

Application Number Priority Date Filing Date Title
KR1020027001080A KR20020019124A (en) 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same
DE60121162T DE60121162T2 (en) 2000-05-26 2001-02-14 COLD-ROLLED STEEL PLATE AND GALVANIZED STEEL PLATE WITH GOOD RECOILING CHARACTERISTICS AND METHOD OF MANUFACTURING THEREOF
CA002379698A CA2379698C (en) 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability
EP01906128A EP1291448B1 (en) 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same
US10/654,775 US7101445B2 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
US10/655,361 US7067023B2 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same

Applications Claiming Priority (8)

Application Number Priority Date Filing Date Title
JP2000/156274 2000-05-26
JP2000156274A JP4524859B2 (en) 2000-05-26 2000-05-26 Cold-drawn steel sheet for deep drawing with excellent strain age hardening characteristics and method for producing the same
JP2000/193717 2000-06-28
JP2000193717 2000-06-28
JP2000/328924 2000-10-27
JP2000328924 2000-10-27
JP2000335803A JP4665302B2 (en) 2000-11-02 2000-11-02 High-tensile cold-rolled steel sheet having high r value, excellent strain age hardening characteristics and non-aging at room temperature, and method for producing the same
JP2000/335803 2000-11-02

Related Child Applications (7)

Application Number Title Priority Date Filing Date
US10/031,426 A-371-Of-International US20030015263A1 (en) 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same
US10031426 A-371-Of-International 2001-02-14
US10/655,394 Division US20040105997A1 (en) 2000-05-26 2003-09-04 Cold-rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
US10/655,288 Division US20040261920A1 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
US10/654,774 Division US20040103962A1 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
US10/654,775 Division US7101445B2 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
US10/655,361 Division US7067023B2 (en) 2000-05-26 2003-09-04 Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same

Publications (1)

Publication Number Publication Date
WO2001090431A1 true WO2001090431A1 (en) 2001-11-29

Family

ID=27481319

Family Applications (1)

Application Number Title Priority Date Filing Date
PCT/JP2001/001004 WO2001090431A1 (en) 2000-05-26 2001-02-14 Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same

Country Status (7)

Country Link
EP (3) EP1498507B1 (en)
KR (1) KR20020019124A (en)
CN (1) CN1158398C (en)
CA (1) CA2379698C (en)
DE (3) DE60121162T2 (en)
TW (1) TW565621B (en)
WO (1) WO2001090431A1 (en)

Cited By (4)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US7559997B2 (en) 2002-06-25 2009-07-14 Jfe Steel Corporation High-strength cold rolled steel sheet and process for producing the same
US8257517B2 (en) 2005-08-03 2012-09-04 Sumitomo Metal Industries, Ltd. Method of producing cold-rolled steel sheet
CN113462963A (en) * 2021-06-10 2021-10-01 江阴市万众精密机械有限公司 Impact-resistant and low-temperature-resistant thrust disc for speed-increasing box coupling and preparation method thereof
WO2022131802A1 (en) * 2020-12-16 2022-06-23 주식회사 포스코 High strength cold rolled, plated steel sheet for home applicances having excellent homogeneous material properties, and method for manufacturing same

Families Citing this family (59)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JP4635525B2 (en) * 2003-09-26 2011-02-23 Ｊｆｅスチール株式会社 High-strength steel sheet excellent in deep drawability and manufacturing method thereof
KR100884104B1 (en) 2004-01-14 2009-02-19 신닛뽄세이테쯔 카부시키카이샤 Hot dip zinc plated high strength steel sheet excellent in plating adhesiveness and hole expanding characteristics
JP4559918B2 (en) * 2004-06-18 2010-10-13 新日本製鐵株式会社 Steel plate for tin and tin free steel excellent in workability and method for producing the same
DE102004044022A1 (en) * 2004-09-09 2006-03-16 Salzgitter Flachstahl Gmbh Well-tempered, unalloyed or microalloyed bake-hardening rolled steel and process for its preparation
US8337643B2 (en) 2004-11-24 2012-12-25 Nucor Corporation Hot rolled dual phase steel sheet
US7442268B2 (en) 2004-11-24 2008-10-28 Nucor Corporation Method of manufacturing cold rolled dual-phase steel sheet
US7959747B2 (en) 2004-11-24 2011-06-14 Nucor Corporation Method of making cold rolled dual phase steel sheet
US11155902B2 (en) 2006-09-27 2021-10-26 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
US7608155B2 (en) 2006-09-27 2009-10-27 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
JP5162924B2 (en) 2007-02-28 2013-03-13 Ｊｆｅスチール株式会社 Steel plate for can and manufacturing method thereof
KR100868457B1 (en) * 2007-05-31 2008-11-11 주식회사 포스코 Galvannealed steel sheet having superior adhesiveness of plated film and method for manufacturing the same
AU2008311043B2 (en) 2007-10-10 2013-02-21 Nucor Corporation Complex metallographic structured steel and method of manufacturing same
US20090236068A1 (en) 2008-03-19 2009-09-24 Nucor Corporation Strip casting apparatus for rapid set and change of casting rolls
CN102015155B (en) 2008-03-19 2013-11-27 纽科尔公司 Strip casting apparatus with casting roll positioning
JP4712882B2 (en) * 2008-07-11 2011-06-29 株式会社神戸製鋼所 High strength cold-rolled steel sheet with excellent hydrogen embrittlement resistance and workability
JP5391607B2 (en) * 2008-08-05 2014-01-15 Ｊｆｅスチール株式会社 High-strength hot-dip galvanized steel sheet with excellent appearance and method for producing the same
JP2010126757A (en) * 2008-11-27 2010-06-10 Jfe Steel Corp High-strength hot-dip galvanized steel sheet and method for producing the same
KR20100076744A (en) * 2008-12-26 2010-07-06 주식회사 포스코 Annealing apparatus of steel sheet, manufacturing apparatus and method for hot-dip galvanized steel with excellent coating quality
JP5051247B2 (en) * 2010-01-15 2012-10-17 Ｊｆｅスチール株式会社 Cold-rolled steel sheet excellent in formability and shape freezing property and its manufacturing method
JP5018935B2 (en) * 2010-06-29 2012-09-05 Ｊｆｅスチール株式会社 High-strength hot-dip galvanized steel sheet excellent in workability and manufacturing method thereof
TWI475114B (en) * 2010-08-18 2015-03-01 Jfe Steel Corp High strength cold rolled steel sheet excellent in workability and impact resistance and a method of manufacturing the same
JP5825481B2 (en) * 2010-11-05 2015-12-02 Ｊｆｅスチール株式会社 High-strength cold-rolled steel sheet excellent in deep drawability and bake hardenability and its manufacturing method
CN102041444A (en) * 2010-12-21 2011-05-04 南阳汉冶特钢有限公司 Low-carbon low-silicon high-quality carbon structural steel and production method thereof
JP5338873B2 (en) 2011-08-05 2013-11-13 Ｊｆｅスチール株式会社 High-strength hot-dip galvanized steel sheet excellent in workability with a tensile strength of 440 MPa or more and its production method
JP2013064169A (en) * 2011-09-15 2013-04-11 Jfe Steel Corp High-strength steel sheet and plated steel sheet excellent in bake-hardenability and formability, and method for production thereof
JP2013072107A (en) * 2011-09-27 2013-04-22 Jfe Steel Corp Bake-hardenable, cold-rolled steel sheet excellent in surface quality after molding and method for manufacturing the same
JP2013072110A (en) * 2011-09-27 2013-04-22 Jfe Steel Corp High-tensile, cold-rolled steel sheet excellent in surface quality after molding and method for manufacturing the same
RU2585889C2 (en) * 2011-09-30 2016-06-10 Ниппон Стил Энд Сумитомо Метал Корпорейшн High-strength hot-dip galvanised steel sheet having excellent resistance to delayed fracture, and method for production thereof
JP2013076132A (en) * 2011-09-30 2013-04-25 Jfe Steel Corp High strength thin steel sheet having excellent bake hardenability and formability and method for manufacturing the same
TWI468530B (en) * 2012-02-13 2015-01-11 新日鐵住金股份有限公司 Cold rolled steel plate, plated steel plate, and method of manufacturing the same
KR101630557B1 (en) * 2012-03-06 2016-06-14 제이에프이 스틸 가부시키가이샤 Warm press forming method and automobile frame component
JP2013209728A (en) * 2012-03-30 2013-10-10 Jfe Steel Corp Cold rolled steel sheet excellent in aging resistance and manufacturing method thereof
JP2013209725A (en) * 2012-03-30 2013-10-10 Jfe Steel Corp Cold rolled steel sheet excellent in bendability and manufacturing method thereof
IN2014DN08471A (en) 2012-04-13 2015-05-08 Nippon Steel & Sumitomo Metal Corp
KR101412293B1 (en) * 2012-04-25 2014-06-25 현대제철 주식회사 High strength cold-rolled steel sheet and method of manufacturing the same
JP2013231216A (en) * 2012-04-27 2013-11-14 Jfe Steel Corp High strength cold rolled steel sheet having excellent chemical conversion property and method for producing the same
JP2014019928A (en) * 2012-07-20 2014-02-03 Jfe Steel Corp High strength cold rolled steel sheet and method for producing high strength cold rolled steel sheet
US20140261916A1 (en) * 2013-03-15 2014-09-18 Thyssenkrupp Steel Usa, Llc High strength - high ductility cold rolled recovery annealed steel and process for manufacture thereof
KR101294575B1 (en) * 2013-03-26 2013-08-07 현대제철 주식회사 Method of manufacturing cold-rolled steel sheet for enamel with excellent fish scale resistance
KR101294477B1 (en) * 2013-03-26 2013-08-07 현대제철 주식회사 Method of manufacturing steel sheet for enamel with excellent dip-drawing property and adhesion
KR101294447B1 (en) * 2013-03-26 2013-08-07 현대제철 주식회사 Method of manufacturing cold-rolled steel sheet for enamel with excellent formability
KR101294577B1 (en) * 2013-03-26 2013-08-07 현대제철 주식회사 Method of manufacturing cold-rolled steel sheet for enamel with excellent dip-drawing property
JP5633594B2 (en) * 2013-04-02 2014-12-03 Ｊｆｅスチール株式会社 Cold-rolled steel sheet excellent in punchability and heat-strain resistance and method for producing the same
US20160201175A1 (en) * 2013-10-02 2016-07-14 Nippon Steel & Sumitomo Metal Corporation Age-hardenable steel
CN103590352B (en) * 2013-11-06 2015-07-08 湖北秭归新亚交通设施有限公司 Processing technology for waste hot-dip galvanized road steel barrier
MY174534A (en) * 2013-12-19 2020-04-24 Nisshin Steel Co Ltd Steel sheet hot-dip coated with zn-al-mg-based system having excellent workability and method for manufacturing same
KR101657845B1 (en) * 2014-12-26 2016-09-20 주식회사 포스코 High strength cold rolled steel sheet having excellent surface quality of thin slab and method for manufacturing the same
CN106244914A (en) * 2015-06-13 2016-12-21 郭策 The preparation method of outdoor monitoring device
CN104950777B (en) * 2015-06-13 2017-11-07 合肥荣事达电子电器集团有限公司 A kind of day and night general outdoor monitoring device
KR102031449B1 (en) * 2017-12-24 2019-10-11 주식회사 포스코 Zinc-based metal plated steel sheet having excellent anti-aging property at room temperature and bake hardenability, and manufacturing method for the same
RU2690076C1 (en) * 2018-12-18 2019-05-30 Публичное акционерное общество "Северсталь" Rolled sheet and method of its production
RU2723872C1 (en) * 2019-05-23 2020-06-17 Публичное акционерное общество "Магнитогорский металлургический комбинат" Method for production of cold-rolled high-strength rolled low-alloy steel sheet
KR102201438B1 (en) * 2019-06-26 2021-01-12 현대제철 주식회사 Steel plate and method of manufacturing the same
KR20220102647A (en) * 2020-01-08 2022-07-20 닛폰세이테츠 가부시키가이샤 Steel plate and its manufacturing method
CN111663075B (en) * 2020-04-09 2021-10-01 北京首钢股份有限公司 Pickled steel for stamping and preparation method and application thereof
CN111926252B (en) * 2020-07-31 2022-01-18 马鞍山钢铁股份有限公司 Hot-rolled pickled steel plate for deep drawing and production method thereof
KR102426248B1 (en) * 2020-11-05 2022-07-28 주식회사 포스코 Method for manufacturing hot-dip galvanized high strength steel sheet having excellent distinctness of image after painting
KR102468036B1 (en) * 2020-11-12 2022-11-17 주식회사 포스코 High-strength hot-dip galvanized steel sheet with excellent formability and process for producing same
TWI768666B (en) * 2021-01-20 2022-06-21 中國鋼鐵股份有限公司 Cold-rolled steel material with high formability and method for producing the same

Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5842753A (en) * 1981-09-07 1983-03-12 Kobe Steel Ltd High gamma value type high strength cold rolled steel plate having composite structure and its manufacture
JPH06116682A (en) * 1992-10-06 1994-04-26 Kawasaki Steel Corp Thin steel sheet for high strength can having baking hardenability and production thereof
EP0608430A1 (en) * 1992-06-22 1994-08-03 Nippon Steel Corporation Cold-rolled steel plate having excellent baking hardenability, non-cold-ageing characteristics and moldability, and molten zinc-plated cold-rolled steel plate and method of manufacturing the same
JPH0734136A (en) * 1993-07-19 1995-02-03 Nippon Steel Corp Production of baking hardenability high strength cold rolled steel sheet excellent in workability
JPH0835039A (en) * 1994-07-21 1996-02-06 Kawasaki Steel Corp Steel sheet for can manufacturing, excellent in baking hardenability and aging resistance and having high strength and high workability, and its production
JPH11279691A (en) * 1998-03-27 1999-10-12 Nippon Steel Corp High strength hot dip galvannealed steel sheet good in workability and its production
JP6102816B2 (en) * 2014-03-31 2017-03-29 ブラザー工業株式会社 Film forming apparatus, film forming method, and film forming program

Family Cites Families (18)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS55141526A (en) * 1979-04-18 1980-11-05 Kawasaki Steel Corp Production of high tension cold-rolled steel plate for deep drawing
JPS5773132A (en) * 1980-10-24 1982-05-07 Nippon Kokan Kk <Nkk> Production of cold rolled mild steel plate of superior deep drawability and aging resistance by continuous annealing
JPS6052528A (en) * 1983-09-02 1985-03-25 Kawasaki Steel Corp Production of high-strength thin steel sheet having good ductility and spot weldability
JPS60145355A (en) * 1984-01-06 1985-07-31 Kawasaki Steel Corp Low yield ratio high tension hot rolled steel sheet having good ductility without deterioration with age and its production
US4578124A (en) * 1984-01-20 1986-03-25 Kabushiki Kaisha Kobe Seiko Sho High strength low carbon steels, steel articles thereof and method for manufacturing the steels
JPS61272323A (en) * 1985-05-28 1986-12-02 Kawasaki Steel Corp Manufacture of original sheet for surface treatment by continuous annealing
NL8502145A (en) * 1985-07-29 1987-02-16 Hoogovens Groep Bv HARD CAN MANUFACTURED FROM A1 QUIET, CONTINUOUS CASTING, CARBON MANGANUM STEEL AND METHOD FOR MANUFACTURING SUCH CAN.
JPH0823048B2 (en) * 1990-07-18 1996-03-06 住友金属工業株式会社 Method for producing hot rolled steel sheet with excellent bake hardenability and workability
EP0475096B2 (en) * 1990-08-17 2004-01-14 JFE Steel Corporation High strength steel sheet adapted for press forming and method of producing the same
US5123969A (en) * 1991-02-01 1992-06-23 China Steel Corp. Ltd. Bake-hardening cold-rolled steel sheet having dual-phase structure and process for manufacturing it
CA2067043C (en) * 1991-04-26 1998-04-28 Susumu Okada High strength cold rolled steel sheet having excellent non-aging property at room temperature and suitable for drawing and method of producing the same
US5690755A (en) * 1992-08-31 1997-11-25 Nippon Steel Corporation Cold-rolled steel sheet and hot-dip galvanized cold-rolled steel sheet having excellent bake hardenability, non-aging properties at room temperature and good formability and process for producing the same
US5486241A (en) * 1992-09-14 1996-01-23 Nippon Steel Corporation Non-aging at room temperature ferritic single-phase cold-rolled steel sheet and hot-dip galvanized steel sheet for deep drawing having excellent fabrication embrittlement resistance and paint-bake hardenability and process for producing the same
JP3458416B2 (en) * 1993-09-21 2003-10-20 Ｊｆｅスチール株式会社 Cold rolled thin steel sheet excellent in impact resistance and method for producing the same
EP0659890B1 (en) * 1993-12-21 2000-03-29 Kawasaki Steel Corporation Method of manufacturing small planar anisotropic high-strength thin can steel plate
CN1152340A (en) * 1995-03-27 1997-06-18 新日本制铁株式会社 Ultralow-carbon cold-rolled sheet and galvanized sheet both excellent in fatigue characteristics and process for producing both
JP4177478B2 (en) * 1998-04-27 2008-11-05 Ｊｆｅスチール株式会社 Cold-rolled steel sheet, hot-dip galvanized steel sheet excellent in formability, panel shape, and dent resistance, and methods for producing them
CN1092714C (en) * 1997-09-04 2002-10-16 川崎制铁株式会社 Steel plates for drum cans, method of manufacturing same, and drum can

Patent Citations (7)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5842753A (en) * 1981-09-07 1983-03-12 Kobe Steel Ltd High gamma value type high strength cold rolled steel plate having composite structure and its manufacture
EP0608430A1 (en) * 1992-06-22 1994-08-03 Nippon Steel Corporation Cold-rolled steel plate having excellent baking hardenability, non-cold-ageing characteristics and moldability, and molten zinc-plated cold-rolled steel plate and method of manufacturing the same
JPH06116682A (en) * 1992-10-06 1994-04-26 Kawasaki Steel Corp Thin steel sheet for high strength can having baking hardenability and production thereof
JPH0734136A (en) * 1993-07-19 1995-02-03 Nippon Steel Corp Production of baking hardenability high strength cold rolled steel sheet excellent in workability
JPH0835039A (en) * 1994-07-21 1996-02-06 Kawasaki Steel Corp Steel sheet for can manufacturing, excellent in baking hardenability and aging resistance and having high strength and high workability, and its production
JPH11279691A (en) * 1998-03-27 1999-10-12 Nippon Steel Corp High strength hot dip galvannealed steel sheet good in workability and its production
JP6102816B2 (en) * 2014-03-31 2017-03-29 ブラザー工業株式会社 Film forming apparatus, film forming method, and film forming program

Non-Patent Citations (1)

* Cited by examiner, † Cited by third party
Title
See also references of EP1291448A4 *

Cited By (5)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US7559997B2 (en) 2002-06-25 2009-07-14 Jfe Steel Corporation High-strength cold rolled steel sheet and process for producing the same
US8257517B2 (en) 2005-08-03 2012-09-04 Sumitomo Metal Industries, Ltd. Method of producing cold-rolled steel sheet
WO2022131802A1 (en) * 2020-12-16 2022-06-23 주식회사 포스코 High strength cold rolled, plated steel sheet for home applicances having excellent homogeneous material properties, and method for manufacturing same
KR20230113515A (en) 2020-12-16 2023-07-31 주식회사 포스코 High strength cold_rolled, plated steel sheet for home applicances having excellent homogeneous material properties, and method for manufacturing the same
CN113462963A (en) * 2021-06-10 2021-10-01 江阴市万众精密机械有限公司 Impact-resistant and low-temperature-resistant thrust disc for speed-increasing box coupling and preparation method thereof

Also Published As

Publication number Publication date
DE60121162T2 (en) 2006-11-09
EP1291448A4 (en) 2004-06-30
EP1291448B1 (en) 2006-06-28
DE60121162D1 (en) 2006-08-10
EP1498507A1 (en) 2005-01-19
CA2379698A1 (en) 2001-11-29
KR20020019124A (en) 2002-03-09
EP1498507B1 (en) 2006-06-28
TW565621B (en) 2003-12-11
CA2379698C (en) 2009-02-17
CN1386140A (en) 2002-12-18
EP1498506A1 (en) 2005-01-19
DE60121233D1 (en) 2006-08-10
DE60121233T2 (en) 2006-11-09
CN1158398C (en) 2004-07-21
DE60121234D1 (en) 2006-08-10
EP1498506B1 (en) 2006-06-28
EP1291448A1 (en) 2003-03-12
DE60121234T2 (en) 2006-11-09

Similar Documents

Publication Publication Date Title
WO2001090431A1 (en) Cold rolled steel sheet and galvanized steel sheet having strain aging hardening property and method for producing the same
JP5365216B2 (en) High-strength steel sheet and its manufacturing method
KR101399741B1 (en) High-strength hot-dip zinc plated steel sheet excellent in workability and process for manufacturing the same
JP4635525B2 (en) High-strength steel sheet excellent in deep drawability and manufacturing method thereof
JP4700764B2 (en) High-strength cold-rolled steel sheet excellent in formability and weldability, high-strength galvanized steel sheet, high-strength galvannealed steel sheet, and methods for producing them
KR101001420B1 (en) Galvanized steel sheet and method for producing the same
KR101264574B1 (en) Method for producing high-strength steel plate having superior deep drawing characteristics
EP2623622B1 (en) High-strength hot-dip galvanized steel sheet with excellent deep drawability and stretch flangeability, and process for producing same
WO2001064967A1 (en) High tensile cold-rolled steel sheet having excellent strain aging hardening properties
WO2001092593A1 (en) Cold-rolled steel sheet having excellent strain aging hardening properties and method for producing the same
JP3846206B2 (en) High tensile cold-rolled steel sheet with excellent strain age hardening characteristics and method for producing the same
KR20150028366A (en) High-strength hot-dip-galvanized steel sheet having excellent moldability, and method for production thereof
JP4501699B2 (en) High-strength steel sheet excellent in deep drawability and stretch flangeability and method for producing the same
JP4206642B2 (en) High tensile hot-rolled steel sheet with excellent strain age hardening characteristics and method for producing the same
JP5251207B2 (en) High strength steel plate with excellent deep drawability and method for producing the same
JP5141232B2 (en) High-strength hot-dip galvanized steel sheet with excellent formability and manufacturing method thereof
KR100933882B1 (en) Manufacturing method of hot dip galvanized steel sheet with excellent workability
JP2002146478A (en) HIGH TENSILE STRENGTH COLD ROLLED STEEL SHEET HAVING HIGH r-VALUE, EXCELLENT STRAIN AGE HARDENING CHARACTERISTIC AND COLD NON-AGING PROPERTY AND ITS PRODUCTION METHOD
JP4556348B2 (en) Ultra-high strength hot-rolled steel sheet with excellent strain age hardening characteristics and method for producing the same
JP2010018852A (en) High strength steel sheet having excellent deep drawability and method for producing the same
JP2002206138A (en) High tensile strength cold rolled steel sheet having excellent formability, strain age hardening characteristic and cold aging resistance and production method therefor
JP4525386B2 (en) Manufacturing method of high-strength steel sheets with excellent shape freezing and deep drawability
JP5655436B2 (en) High-strength steel sheet excellent in deep drawability and manufacturing method thereof
JP5440370B2 (en) Alloyed hot-dip galvanized steel sheet and method for producing the same
JP4517629B2 (en) Composite structure cold-rolled steel sheet, plated steel sheet having excellent surface strain resistance, and production method thereof

Legal Events

Date Code Title Description
AK Designated states

Kind code of ref document: A1

Designated state(s): CA CN KR US

AL Designated countries for regional patents

Kind code of ref document: A1

Designated state(s): AT BE CH CY DE DK ES FI FR GB GR IE IT LU MC NL PT SE TR

WWE Wipo information: entry into national phase

Ref document number: 2001906128

Country of ref document: EP

Ref document number: 2379698

Country of ref document: CA

WWE Wipo information: entry into national phase

Ref document number: 10031426

Country of ref document: US

121 Ep: the epo has been informed by wipo that ep was designated in this application
WWE Wipo information: entry into national phase

Ref document number: 1020027001080

Country of ref document: KR

WWP Wipo information: published in national office

Ref document number: 1020027001080

Country of ref document: KR

WWE Wipo information: entry into national phase

Ref document number: 018021867

Country of ref document: CN

WWP Wipo information: published in national office

Ref document number: 2001906128

Country of ref document: EP

WWR Wipo information: refused in national office

Ref document number: 1020027001080

Country of ref document: KR

WWG Wipo information: grant in national office

Ref document number: 2001906128

Country of ref document: EP