EP1291448A1 - Kaltgewalztes stahlblech und galvanisiertes stahlblech mit guten reckalterungseigenschaften und herstellungsverfahren dafür - Google Patents

Kaltgewalztes stahlblech und galvanisiertes stahlblech mit guten reckalterungseigenschaften und herstellungsverfahren dafür Download PDF

Info

Publication number
EP1291448A1
EP1291448A1 EP01906128A EP01906128A EP1291448A1 EP 1291448 A1 EP1291448 A1 EP 1291448A1 EP 01906128 A EP01906128 A EP 01906128A EP 01906128 A EP01906128 A EP 01906128A EP 1291448 A1 EP1291448 A1 EP 1291448A1
Authority
EP
European Patent Office
Prior art keywords
less
cold
rolled
steel sheet
sheet
Prior art date
Legal status (The legal status 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 status listed.)
Granted
Application number
EP01906128A
Other languages
English (en)
French (fr)
Other versions
EP1291448A4 (de
EP1291448B1 (de
Inventor
Chikara Technical Research Laboratories Kami
Akio Chiba Works Tosaka
Takuya Technical Research Laboratories Yamazaki
Current Assignee (The listed assignees may be inaccurate. Google has not performed a legal analysis and makes no representation or warranty as to the accuracy of the list.)
JFE Steel Corp
Original Assignee
JFE Steel Corp
Kawasaki Steel Corp
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/ja
Priority claimed from JP2000335803A external-priority patent/JP4665302B2/ja
Application filed by JFE Steel Corp, Kawasaki Steel Corp filed Critical JFE Steel Corp
Priority to EP04023101A priority Critical patent/EP1498507B1/de
Priority to EP04023082A priority patent/EP1498506B1/de
Publication of EP1291448A1 publication Critical patent/EP1291448A1/de
Publication of EP1291448A4 publication Critical patent/EP1291448A4/de
Application granted granted Critical
Publication of EP1291448B1 publication Critical patent/EP1291448B1/de
Anticipated expiration legal-status Critical
Expired - Lifetime legal-status Critical Current

Links

Images

Classifications

    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—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
    • 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
    • C21D8/0221—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 relates to a cold-rolled steel sheet, an electro-galvanized steel sheet, a hot-dip galvanized steel sheet, and an alloyed hot-dip galvanized steel sheet, which are suitable as raw material steel sheets for molded products such as building members, mechanical structural parts, automobile structural parts, etc., which are used at positions required to have structural strength, particularly, strength and/or stiffness in deformation, and which are subjected to heat treatment for increasing strength after processing such as pressing or the like, and a method of producing these steel sheets.
  • a process of coating and baking at lower than 200°C is used as a method in which a material having low deformation stress before press forming to facilitate press forming, and then hardened after press forming to increase the strength of a part.
  • a steel sheet for such coating and baking a BH steel sheet has been developed.
  • Japanese Unexamined Patent Application Publication No. 55-141526 discloses a method in which Nb is added according to the contents of C, N and Al of steel, Nb/(dissolved C + dissolved N) by at% is limited in a specified range, and the cooling rate after annealing is controlled to adjust dissolved C and dissolved N in a steel sheet.
  • Japanese Examined Patent Application Publication No. 61-45689 discloses a method in which baking hardenability is improved by adding Ti and Nb.
  • Japanese Unexamined Patent Application Publication No. 5-25549 discloses a method in which baking hardenability is improved by adding W, Cr and Mo to steel singly or in a combination.
  • Japanese Unexamined Patent Application Publication No. 10-310847 discloses an alloying ho-dip galvanized steel sheet having tensile strength increased by 60 MPa or more by heat treatment in the temperature region of 200 to 450°C.
  • This steel sheet contains, by mass%, 0.01 to 0.08% of C, and 0.01 to 3.0% of Mn, and at least one of W, Cr, and Mo in a total of 0.05 to 3.0%, and further contains at least one of 0.005 to 0.1% of Ti, 0.005 to 0.1% of Nb and 0.005 to 0.1% of V according to demand, and the microstructure of the steel is composed of ferrite or mainly composed of ferrite.
  • this technique comprises forming a fine carbide in the steel sheet by heat treatment after forming to effectively propagate a dislocation of stress applied during pressing, thereby increasing the amount of strain. Therefore, heat treatment must be performed in the temperature range of 220 to 370°C. There is thus the problem of a necessary heat treatment temperature higher than general bake-hardening temperatures.
  • An automobile part using a high-tensile-strength thin steel sheet must exhibit a sufficient property according to its function.
  • the property depends upon the part, and examples of the property include dent resistance, static strength against deformation such as bending, twisting, or the like, fatigue resistance, impact resistance, etc.
  • the high-tensile-strength steel sheet used for an automobile part is required to be excellent in such a property after forming.
  • the properties are related to the. strength of a steel sheet after forming, and thus the lower limit of strength of the high-tensile-strength steel sheet used must be set for achieving thinning.
  • a steel sheet is press-molded. If the steel sheet has excessively high strength in press forming, the steel sheet causes the following problems: (1) deteriorating shape fixability; (2) deteriorating ductility to cause cracking, necking, or the like during forming; and (3) deteriorating dent resistance (resistance to a dent produced by a local compressive load) when the sheet thickness is decreased. These problems thus inhibit the extension of application of the high-tensile-strength steel sheet to automobile bodies.
  • a steel sheet composed of ultra-low-carbon steel is known as a raw material, for example, for a cold-rolled steel sheet for an external sheet panel, in which the content of C finally remaining in a solid solution state is controlled to an appropriate range.
  • This type of steel sheet is kept soft during press forming to ensure shape fixability and ductility, and its yield stress is increased by utilizing the strain aging phenomenon which occurs in the step of coating and baking at 170°C for about 20 minutes after press forming, to ensure dent resistance.
  • This steel sheet is soft during press forming because C is dissolved in steel, while dissolved C is fixed to a dislocation introduced in press forming in the coating and baking step after press forming to increase the yield stress.
  • Japanese Unexamined Patent Application Publication No. 60-52528 discloses a method of producing a high-strength steel thin sheet having good ductility and spot weldability, in which steel containing 0.02 to 0.15% of C, 0.8 to 3.5% of Mn, 0.02 to 0.15% of P, 0.10% or less of Al, and 0.005 to 0.025% of N is hot-rolled by coiling at a temperature of 550°C or less, cold-rolled, and then annealed by controlled cooling and heat treatment.
  • a steel sheet produced by the technique disclosed in Japanese Unexamined Patent Application Publication No. 60-52528 has a mixed structure comprising a low-temperature transformation product phase mainly composed of ferrite and martensite, and having excellent ductility, and high strength is achieved by utilizing strain aging due to positively added N during coating baking.
  • Japanese Examined Patent Application Publication No. 5-24979 discloses a high-tensile-strength cold-rolled steel thin sheet having baking hardenability which has a composition comprising 0.08 to 0.20% of C, 1.5 to 3.5% of Mn, and the balance composed of Fe and inevitable impurities, and a structure composed of homogeneous bainite containing 5% or less of ferrite, or bainite partially containing martensite.
  • 5-24979 is produced by quenching in the temperature range of 200 to 400°C in the cooling process after continuous annealing, and then slowly cooling to obtain a structure mainly composed of bainite and having a large amount of bake-hardening which is not obtained by a conventional method.
  • Japanese Examined Patent Application Publication No. 61-12008 discloses a method of producing a high-tensile-strength steel sheet having a high r value.
  • This method is characterized by annealing ultra-low-C steel used as a raw material in a ferrite-austenite coexistence region after cold rolling.
  • the resultant steel sheet has a high r value and a high degree of baking hardenability (BH property), but the obtained BH amount is about 60 MPa at most.
  • the yield point of the steel sheet is increased after strain aging, but TS is not increased, thereby causing the problem of limiting application to parts.
  • the above-described steel sheet exhibits excellent strength after coating and baking in a simple tensile test, but produces large variations in strength during plastic deformation under actual pressing conditions. Therefore, it cannot be said that the steel sheet is sufficiently applied to parts required to have reliability.
  • Japanese Examined Patent Application Publication No. 8-23048 discloses a method of producing a hot-rolled steel sheet which is soft during processing, and has tensile strength increased by coating and baking after processing to be effective to improve fatigue resistance.
  • steel contains 0.02 to 0.13 mass % of C, and 0.0080 to 0.0250 mass % of N, and the finisher deliver temperature and the coiling temperature are controlled to leave a large amount of dissolved N in the steel, thereby forming a composite structure as a metal structure mainly composed of ferrite and martensite. Therefore, an increase of 100 MPa or more in tensile strength is achieved at the heat treatment temperature of 170°C after forming.
  • Japanese Unexamined Patent Application Publication No. 10-183301 discloses a hot-rolled steel sheet having excellent baking hardenability and natural aging resistance, in which the C and N contents are limited to 0.01 to 0.12 mass % and 0.0001 to 0.01 mass %, respectively, and the average crystal grain diameter is controlled to 8 ⁇ m or less to ensure a BH amount of as high as 80 MPa or more, and suppress the AI amount to 45 MPa or less.
  • this steel sheet is a hot-rolled sheet, and is thus difficult to obtain a high r value because the ferrite aggregation texture is made random due to auste141-ferrite transformation. Therefore, the steel sheet cannot be said to have sufficient deep drawability.
  • the hot-rolled steel sheet obtained by this technique is used as a starting material for cold rolling and recrystallization annealing, the increase in tensile strength obtained after forming and heat treatment is not always equivalent to a hot-rolled steel sheet, and a BH amount of as high as 80 MPa or more cannot be always obtained.
  • the microstructure of the cold -rolled steel becomes different from that of hot-rolled one due to cold rolling and recrystallization annealing, and strain greatly accumulates during cold rolling to easily form a carbide, a nitride or a carbonitride, thereby changing the states of dissolved C and dissolved N.
  • an object of the present invention is to provide a cold-rolled steel sheet and a hot-dip galvanized steel sheet (including an alloyed steel sheet) for deep drawing, which have excellent deep drawability, TS x r value ⁇ 750 MPa, and excellent strain aging hardenability (BH ⁇ 80 MPa and ⁇ TS ⁇ 40 MPa), and an effective method of producing these steel sheets.
  • Another object of the present invention is to solve the above problems of the conventional techniques and provide a high-tensile-strength cold-rolled steel sheet which is suitable for automobile parts required to have high moldability, softness and high moldability, and stable material properties, and which can easily be molded to an automobile part having a complicated shape without producing shape defects such as spring back, twisting, and curving, and cracking, etc., and which has sufficient strength as an automobile part after heat treatment of a molded automobile part to permit sufficient contribution to a reduction in body weight of an automobile, a high r value of 1.2 or more, and excellent strain age hardenability.
  • a further object of the present invention is to provide an industrial production method capable of producing the steel sheet at low cost without disturbing its shape.
  • the inventors produced various steel sheets having different compositions under various production conditions, and experimentally evaluated various material properties. As a result, it was found that both moldability and hardenability after forming can be improved by using as a strengthening element N, which has not be positively used before in a field requiring high processability, and effectively using the great strain age hardening phenomenon manifested by the action of the strengthening element.
  • the strain age hardening phenomenon due to N in order to advantageously use the strain age hardening phenomenon due to N, the strain age hardening phenomenon due to N must be advantageously combined with a condition for coating and baking an automobile, or further positively combined with a heat treatment condition after forming. It was thus found to be effective to appropriately control the hot rolling condition, the cold rolling and the cold rolling annealing condition to control the microstructure of a steel sheet and the amount of dissolved N in certain ranges. It was also found that in order to stably manifest the strain age hardening phenomenon due to N, it is important to control the Al content of the composition according to the N content.
  • the inventors further found that in order to obtain a high r value, the C content is decreased, continuous annealing is performed in the ferrite-austenite two-phase temperature region, and subsequent cooling is controlled to form a structure containing an acicular ferrite phase at an area ratio of 5% or more in the ferrite phase.
  • Such a combination of the microstructure and the appropriate amount of dissolved N was found to enable the achievement of a cold-rolled steel sheet having a high r value, excellent press moldability, and excellent strain age hardenability. This was also found to permit sufficient use of N without causing the problem of natural aging deterioration, which is the problem of a conventional bake-hardening steel sheet.
  • the inventors found that by suing N as a strengthening element, controlling the Al content according to the N content in an appropriate range, and appropriately controlling the hot rolling condition and the cold rolling annealing condition to appropriately control the microstructure and dissolved N, it is possible to obtain a steel sheet having a high r value and excellent moldability as compared with conventional solid-solution strengthening-type C-Mn steel sheets and precipitation strengthening-type steel sheets, and strain age hardenability, which is not possessed by the conventional steel sheets.
  • a steel sheet of the present invention exhibits higher strength after coating and baking in a simple tensile test, as compared with a conventional steel sheet, and exhibits small variations in strength in plastic deformation under actual pressing conditions and stable part strength, thereby enabling application to parts required to have reliability.
  • a portion where large strain is applied to decrease the thickness has higher hardenability than other portions, and is considered homogeneous when being evaluated based on a surcharge load ability of (thickness) x (strength), thereby stabilizing strength as a part.
  • the present invention has been achieved based on the above findings.
  • the findings were obtained from the experiment described below.
  • a sheet bar (thickness: 30 mm) having a composition containing, by mass %, 0.0015% of C, 0.0010% of B, 0.015 of Si, 0.5% of Mn, 0.03% of P, 0.08% of S and 0.011% of N, 0.005 to 0.05% of Nb and 0.005 to 0.03% of Al, and the balance composed of Fe and inevitable impurities was uniformly heated at 1150°C, hot-rolled by three passes so that the temperature of the final pass was 900°C higher than the Ar 3 transformation point, and then cooled with water for 0.1 second. Then, the sheet bar was subjected to heat treatment corresponding to coiling at 500°C for 1 hour.
  • the thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealed at 800°C for 40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling direction, and tensile strength was measured with a strain rate of 0.02/s by using a general tensile testing machine. Also, tensile strain of 10% was applied to a tensile test specimen of JIS No.
  • Fig. 1 shows the results of measurement of the relation between the steel compositions (N% - 14/93•Nb% - 14/27•Al%-14/11•B%) and ⁇ TS.
  • ⁇ TS becomes 60 MPa or more when the value of (N% - 14/93•Nb% - 14/27•Al% - 14/11•B%) satisfies 0.0015 mass %.
  • a sheet bar (thickness: 30 mm) having a composition containing, by mass %, 0.0010% of C, 0.02 of Si, 0.6% of Mn, 0.01% of P, 0.009% of S and 0.012% of N, 0.01% of Al, 0.015% of Nb, 0.00005 to 0.0025% of B, and the balance composed of Fe and inevitable impurities was uniformly heated at 1100°C, hot-rolled by three passes so that the temperature of the final pass was 920°C higher than the Ar 3 transformation point, and then cooled with water for 0.1 second. Then, the sheet bar was subjected to heat treatment corresponding to coiling at 450°C for 1 hour.
  • the thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealed at 820°C for 40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%. Then, a tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling direction, and tensile strength was measured with a strain rate of 0.02/s by using a general tensile testing machine. Also, tensile strain of 10% was applied to a tensile test specimen of JIS No. 5 separately obtained from the cold-rolled sheet in the rolling direction, and then the specimen was subjected to a normal tensile test after heat treatment at 120°C for 20 minutes.
  • Fig. 2 shows the results of measurement of the relation between the B content of steel and ⁇ TS. This figure indicates that with a B content of 0.0005 to 0.0015 mass %, a high ⁇ TS of 60 MPa or more can be obtained.
  • steel B having a composition containing, by mass %, 0.010% of C, 0.0012% of N, 0.0010% of B, 0.01% of Si, 0.5% of Mn, 0.03% of P, 0.008% of S, 0.014% of Nb, 0.01% of Al, and the balance composed of Fe and inevitable impurities was uniformly heated at 1150°C, hot-rolled by three passes so that the temperature of the final pass was 910°C higher than the Ar 3 transformation point, and then cooled with a gas for 0.1 second. Then, each of the sheet bars was subjected to heat treatment corresponding to coiling at 600°C for
  • Each of the thus-obtained hot-rolled sheets having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealed at 880°C for 40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%.
  • a tensile test specimen of JIS No. 5 was obtained from each of the resultant cold-rolled sheets in the rolling direction, and tensile strength was measured with a strain rate of 0.02/s by using a general tensile testing machine. Also, tensile strain of 10% was applied to a tensile test specimen of JIS No. 5 separately obtained from each of the cold-rolled sheets in the rolling direction, and then the specimen was subjected to a normal tensile test after heat treatment at various temperatures for 20 minutes.
  • Fig. 3 shows the results of measurement of the influence of the heat treatment temperature after forming on ⁇ TS. This figure indicates that in the relatively low temperature region of heat treatment temperatures of 200°C or less after forming, the ultra-low carbon steel A having a high N content exhibits higher ⁇ TS than the semi-ultra low carbon steel B having a low N content, and while in the high temperature region, both steel materials exhibit substantially the same ⁇ TS. There experimental results reveal that in order to ensure ⁇ TS in the low temperature region, it is effective to use dissolved N.
  • Fig. 4 shows the results of measurement of the influences of the crystal grain diameter d and steel compositions (N%-14/93•Nb% - 14/27•Al% - 14/11•B%) on a decrease ( ⁇ E1) in elongation by natural aging and an increase in tensile strength ( ⁇ TS) after forming.
  • the decrease ( ⁇ E1) in elongation was evaluated by the difference between the total elongation measured with the test specimen of JIS NO. 5 obtained from each of the cold-rolled sheets in the rolling direction, and the total elongation measured with the separately obtained test specimen after holding at 100°C for 8 hours for accelerating natural aging.
  • Fig. 4 indicates that when the value of (N% - 14/93•Nb%-14/27•A1% - 14/11•B%) is 0.0015 mass % or more, and the crystal grain diameter d is 20 ⁇ m or less, both high ⁇ TS and low ⁇ E1 can be achieved.
  • a sheet bar of steel containing 0.0015% of C, 0.30 of Si, 0.8% of Mn, 0.03% of P, 0.005% of S and 0.012% of N, and 0.02 to 0.08% of Al was uniformly heated at 1050°C, hot-rolled by seven passes so that the temperature of the final pass was 670°C, and then recrystallized and annealed at 700°C for 5 hours.
  • the thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealed at 875°C for 40 seconds, and then temper-rolled with a rolling reduction of 0.8%. Then, a tensile test specimen of JIS No.
  • a sheet bar of steel containing 0.0015% of C, 0.0010% of B, 0.01 of Si, 0.5% of Mn, 0.03% of P, 0.008% of S and 0.011% of N, 0.005 to 0.05% of Nb, and 0.005 to 0.03% of Al was uniformly heated at 1000°C, hot-rolled by seven passes so that the temperature of the final pass was 650°C, and then recrystallized and annealed at 800°C for 60 seconds.
  • the thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealing at 880°C for 40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%.
  • a tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling direction, and TS x r value, BH and ⁇ TS were measured with a strain rate of 3 x 10 - 3 /s by using a general tensile testing machine.
  • the relations between the measured values and N/(Al+Nb+B) are shown in Fig. 5.
  • steel containing 0.005 to 0.05% of Nb and 0.0010% of B was used. This figure indicates that in the range of N/(Al+Nb+B) ⁇ 0.30, BH ⁇ 80 MPa, ⁇ TS ⁇ 60 MPa, and TS x r value ⁇ 850 are achieved.
  • a sheet bar of steel containing 0.0010% of C, 0.02 of Si, 0.6% of Mn, 0.01% of P, 0.009% of S and 0.015% of N, 0.01% of Al, 0.015% of Nb and 0.0001 to 0.0025% of B was uniformly heated at 1050°C, hot-rolled by seven passes so that the temperature of the final pass was 680°C, and then recrystallized and annealed at 850°C for 5 hours.
  • the thus-obtained hot-rolled sheet having a thickness of 4 mm was cold-rolled with a rolling reduction ratio of 82,5%, recrystallized and annealed at 880°C for 40 seconds, and then temper-rolled with a rolling reduction ratio of 0.8%.
  • a tensile test specimen of JIS No. 5 was obtained from the resultant cold-rolled sheet in the rolling direction, and TS x r value, BH and ⁇ TS were measured with a strain rate of 3 x 10 -3 /s by using a general tensile testing machine.
  • the relations between the measured values and the B content are shown in Fig. 6.
  • a cold-rolled steel sheet having excellent strain age hardenability comprises a composition, by mass %, comprising:
  • a cold-rolled steel sheet having excellent strain age hardenability comprises a composition, by mass %, comprising:
  • the composition by mass %, preferably further comprises:
  • the above composition by mass %, preferably further comprises at least one of Cu, Ni and Mo in a total of 1.0% or less according to demand.
  • the steel sheet preferably has a crystal grain diameter of 20 ⁇ m or less.
  • strength after forming is preferably increased by 60 MPa or more by heat treatment in the low temperature region of 120 to 200°C.
  • the surface of the cold-rolled steel sheet may be coated by electro-galvanization, hot-dip galvanization, or alloying hot-dip galvanization.
  • a method of producing a cold-rolled steel sheet having excellent strain age hardenability comprises hot-rolling a steel slab under conditions in which the slab is cooled immediately after the end of finish rolling and coiled at a coiling temperature of 400 to 800°C, cold-rolling the hot-rolled sheet with a rolling reduction ratio of 60 to 95%, and then performing recrystallization annealing at a temperature of 650 to 900°C, wherein the steel slab has a composition, by mass %, comprising:
  • the composition by mass %, preferably further comprises:
  • the temperature in the heating-up step in recrystallization annealing, is preferably increased at a rate of 1 to 20°C/s in the temperature region from 500°C to the recrystallization temperature.
  • hot-dip galvanization and heat alloying may be performed after the recrystallization annealing.
  • a cold-rolled deep drawing steel sheet having excellent strain age hardenability comprises a composition, by mass %, comprising:
  • the composition by mass %, preferably further comprises:
  • composition by mass %, preferably further comprises at least one of the following:
  • a method of producing a cold-rolled deep drawing steel sheet having excellent strain age hardenability comprises heating a steel raw material to 950°C or more, roughly rolling the raw material so that the finisher delivery temperature is Ar 3 to 1000°C, finish-rolling the material while lubricating it in the temperature region of 600°C to Ar 3 , coiling the rolled sheet in which the total reduction ratio by rolling starting from rough rolling to finish rolling is 80% or more, recrystallizing and annealing the hot-rolled sheet, cold-rolling the rolled sheet with a rolling reduction ratio of 60 to 95%, and then recrystallizing and annealing the resultant cold-rolled sheet, wherein the steel raw material has a composition, by mass %, comprising:
  • a high-tensile-strength cold-rolled steel sheet having excellent moldability, strain age hardenability and natural aging resistance comprises a composition, by mass %, comprising:
  • composition preferably further comprises at least one of the following groups a to c:
  • a method of producing a high-tensile-strength cold-rolled steel sheet having a r value of 1.2 or more, and excellent moldability, strain age hardenability and natural aging resistance comprises:
  • the composition preferably further comprises, by mass %, at least one of the following groups a to c:
  • a high-tensile-strength cold-rolled steel sheet having a high r value and excellent strain age hardenability and natural aging resistance comprises a composition, by mass %, comprising:
  • composition preferably further comprises at least one of the following groups d to g:
  • a method of producing a high-tensile-strength cold-rolled steel sheet having a r value of as high as 1.2 or more, and excellent strain age hardenability and natural aging resistance comprises:
  • over aging after the cooling step of the continuous annealing is preferably carried out in the temperature range of 350°C or more for a residual time of 20 seconds or more.
  • the composition preferably further comprises, by mass %, at least one of the following groups d to g:
  • C is advantageously as small as possible. Also, redissolution of NbC proceeds in the annealing step after cold rolling to increase the amount of dissolved C in crystal grains, thereby easily causing deterioration in natural aging resistance. Therefore, the C amount is preferably suppressed to less than 0.01 mass %, more preferably 0.0050 mass % or less, and most preferably 0.0030 mass % or less. Si: 0.005 to 1.0 mass %
  • Si is a useful composition for suppressing a decrease in elongation, and improving strength.
  • Si content is limited to the range of 0.005 to 1.0 mass %, and preferably the range of 0.01 to 0.75 mass %.
  • Mn 0.01 to 1.5 mass %
  • Mn not only is useful as a strengthening composition for steel, but also has the function to suppress embrittlement with S due to the formation of MnS.
  • Mn content is limited to the range of 0.01 to 1.5 mass %, and preferably the range of 1.10 to 0.75 mass %.
  • P is a solid solution strengthening element which effectively contributes to reinforcement of steel.
  • P content of over 0.01 mass %, deep drawability deteriorates due to the formation of phosphide such as (FeNb) x P or the like. Therefore, P is limited to 0.10 mass % or less.
  • Al is added as a deoxidizer for improving the yield of carbonitride forming components.
  • Al content is less than 0.005 mass %, the effect is insufficient, while with an Al content of over 0.030 mass %, the amount of N to be added to steel is increased to easily cause slab defects during steel making. Therefore, Al is contained in the range of 0.005 to 0.030 mass %.
  • N is an important element which plays the role of imparting strain age hardenability to a steel sheet.
  • N content is contained in the range of 0.005 to 0.040 mass %, and preferably in the range of 0.008 to 0.015 mass %.
  • B is added in a combination with Nb to exhibit the function to effectively make fine the hot-rolled structure and the cold-rolled recrystallized annealed structure and to improve cold-work embrittlement resistance.
  • B content is contained in the range of 0.0001 to 0.003 mass %, preferably in the range of 0.0001 to 0.0015 mass %, and more preferably in the range of 0.0007 to 0.0012 mass %.
  • Nb 0.005 to 0.050 mass %
  • Nb is added in a combination with B to contribute to refinement of the hot-rolled structure and the cold-rolled recrystallized annealed structure, and have the function to fix dissolved C as NbC. Furthermore, Nb forms a nitride NbN to contribute to refinement of the cold-rolled recrystallized annealed structure.
  • Nb is contained in the range of 0.005 to 0.050 mass %, and preferably 0.010 to 0.030 mass %.
  • Nb has the function to fix dissolved C as NbC, and forms a nitride NbN.
  • Al and B form A1N and BN, respectively. Therefore, in order to ensure the sufficient amount of dissolved N and sufficiently decrease the amount of dissolved C, it is important to satisfy the following relations (1) and (2): N% ⁇ 0.0015 + 14/93•Nb% + 14/27•Al% + 14/11•B% C% ⁇ 0.5•(12/93)•Nb%
  • the crystal grain diameter is preferably decreased.
  • ⁇ E1 can be suppressed to 2.0% or less by decreasing the crystal grain diameter d to 20 ⁇ m or less.
  • the crystal grain diameter d is more preferably decreased to 15 ⁇ m or less. This is because, as shown in Fig. 4, ⁇ E1 can be suppressed to 2.0% or less by decreasing the crystal grain diameter d to 15 ⁇ m or less.
  • Steel having the above-described suitable composition is melted by a known melting method such as a converter or the like, and a steel slab is formed by an ingot making method or a continuous casting method.
  • the heating temperature of hot rolling is not specified, but the heating temperature of hot rolling is preferably set to 1300°C or less. This is because it is advantageous to fix and precipitate dissolved C as a carbide in order to improve deep drawability.
  • the heating temperature is preferably set to 1150°C or less.
  • the lower limit of the heating temperature is preferably 900°C.
  • the total rolling reduction ratio of hot rolling is preferably 70% or more. This is because with a toal rolling reduction ratio of less than 70%, the crystal grains of the hot-rolled sheet are not sufficiently made fine.
  • finish rolling is preferably finished in the temperature region of 650 to 960°C, and the finishing temperature of hot-rolling may be in the ⁇ region above the Ar 3 transformation point, or the ⁇ region below the Ar 3 transformation pint.
  • the finishing temperature in hot-rolling process over 960°C, the crystal grains of the hot-rolled sheet are coarsened to deteriorate deep drawability after cold rolling and annealing.
  • a temperature of less than 650°C deformation resistance is increased to increase the hot-rolling load, causing difficulties in rolling.
  • cooling is started immediately after the end of final rolling in hot-rolling process to prevent normal grain growth and suppress AlN precipitation in the cooling step.
  • the starting time of the cooling step is preferably within 1.5 seconds, more preferably 1.0 second, and most preferably 0.5 second, after the end of finish rolling. This is because when cooling is performed immediately after the end of rolling, a large amount of ferrite nuclei is produced due to an increase in the degree of over cooling with accumulated strain to promote ferrite transformation and suppress the diffusion of dissolved N in the ⁇ phase into the ferrite grains, thereby increasing the amount of dissolved N present in the ferrite grain boundaries.
  • the cooling rate is preferably 10°C/s or more in order to ensure dissolved N.
  • the cooling rate is preferably 50°C/s or more in order to ensure dissolved N.
  • the hot-rolled sheet is coiled.
  • the coiling temperature is advantageously as high as possible.
  • the scale formed on the surface of the hot-rolled sheet is thickened to increase the load of the work of removing the scale, and progress the formation of a nitride, causing a change in the amount of dissolved N in the coil length direction.
  • the coiling temperature of the hot-rolled sheet must be in the range of 400 to 800°C.
  • the hot-rolled sheet is cold-rolled, but the rolling reduction ratio of cold rolling must be 60 to 95%. This is because with a rolling reduction ratio of cold rolling of less than 60%, a high r value cannot be expected, while with a rolling reduction ratio of over 95%, the r value is decreased.
  • the cold-rolled sheet subjected to cold rolling is then recrystallized and annealed.
  • the annealing method may be either continuous annealing or batch annealing, continuous annealing is advantageous.
  • the continuous annealing may be performed either in a normal continuous annealing line or in a continuous hot-dip galvanization line.
  • the preferable annealing conditions include 650°C or more for 5 seconds or more. This is because with an annealing temperature of less than 650°C, and an annealing condition of less than 5 seconds, recrystallization is not completed to decrease deep drawability. In order to improve deep drawability, annealing is preferably performed in the ferrite single phase region at 800°C or more for 5 seconds or more.
  • Annealing in the high-temperature ⁇ + ⁇ two-phase region partially produces ⁇ ⁇ ⁇ transformation to improve the r value due to the development of the ⁇ 111 ⁇ aggregation structure.
  • the aggregation structure is made random to decrease the r value, thereby deteriorating deep drawability.
  • the upper limit of the annealing temperature is preferably 900°C. This is because with an annealing temperature of over 900°C, redissolution of a carbide proceeds to excessively increase the amount of dissolved C, thereby deteriorating the natural aging property.
  • ⁇ ⁇ ⁇ transformation occurs, the aggregation structure is made random to decrease the r value, deteriorating deep drawability.
  • the temperature region in which controlled heating must be performed is 500°C, at which AlN or the like starts to precipitate, to the recrystallization temperature
  • the heating rate is preferably in the range of 1 to 20°C/s because with a heating rate of over 20°C/s, the sufficient amount of precipitates cannot be obtained, while with a heating rate of less than 1°C/s, precipitates are coarsened to weaken the effect of suppressing grain growth.
  • temper rolling of 10% or less may be performed for correcting the shape and controlling surface roughness.
  • the cooling rate after soaking in recrystallization annealing is preferably 10 to 50°C/s. This is because with a cooling rate of 10°C/s or less, grains are grown during cooling to coarsen the crystal grains, thereby deteriorating the strain aging property and natural aging property. While with a cooling rate of 50°C/s or more, dissolved N does not sufficiently diffuse into the grain boundaries, deteriorating the natural aging property.
  • the cooling rate is preferably 10 to 30°C/s.
  • the hot-dip galvanization and alloying are not limited, and may be performed according to a conventional known method.
  • a steel sheet subjected to surface treatment generally used for steel thin sheets such as a steel sheet (a dull-finished steel sheet, a bright-finished steel sheet, or a steel sheet having a specified roughness pattern formed on the surface thereof), which is produced by temper-rolling the alloyed hot-dip galvanized steel sheet, for improving processability and the appearance after processing, a steel sheet having an oil film layer of antirust oil or lubricating oil formed on the surface thereof, or the like, the effect of the present invention can be sufficiently exhibited in the composition range of the prevent invention.
  • a cold-rolled steel sheet and a galvannealed steel sheet can be obtained, which have excellent deep drawability and excellent strain age hardenability, that tensile strength increased by press forming and heat treatment.
  • C is advantageously as small as possible. Also, redissolution of NbC proceeds in the annealing step after cold rolling to increase the amount of dissolved C in crystal grains, thereby easily causing deterioration in natural aging resistance. Therefore, the C amount is preferably suppressed to less than 0.01 mass %, more preferably 0.0050 mass % or less, and most preferably 0.0030 mass % or less. In order to ensure strength and prevent coarsening of crystal grains, the C content is preferably 0.0005% or more. Si: 0.005 to 1.0 mass %
  • Si is a useful component for suppressing a decrease in elongation, and improving strength.
  • the Si content is limited to the range of 0.005 to 1.0 mass %, and preferably the range of 0.01 to 0.75 mass %.
  • Mn 0.01 to 1.5 mass %
  • Mn not only is useful as a strengthening component for steel, but also has the function to suppress embrittlement with S due to the formation of MnS..
  • Mn content is limited to the range of 0.01 to 1.5 mass %, and preferably the range of 0.10 to 0.75 mass %.
  • P is a solid solution strengthening element which effectively contributes to strengthening of steel.
  • P content of over 0.01 mass %, deep drawability deteriorates due to the formation of phosphide such as (FeNb) x P or the like. Therefore, P is limited to 0.10 mass % or less. S: 0.01 mass % or less
  • Al is added as a element for deoxidization for improving the yield of the elements forming carbonitride.
  • Al content is contained in the range of 0.005 to 0.030 mass %.
  • N is an important element which plays the role of imparting strain age hardenability to a steel sheet.
  • N content is contained in the range of 0.005 to 0.040 mass %, and preferably in the range of 0.008 to 0.015 mass %.
  • B is added in a combination with Nb to exhibit the function to effectively make fine the micro structure of the hot-rolled steel and the cold-rolled steel, annealed for recrystallization, and to improve cold-work embrittlement resistance.
  • B is contained in the range of 0.0001 to 0.003 mass %, preferably in the range of 0.0001 to 0.0015 mass %, and more preferably in the range of 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 are added in a combination with B to contribute to refinement of the the micro structure of the hot-rolled steel and the cold-rolled steel, annealed for recrystallization, and have the function to precipitate dissolved C as NbC, Tic and VC, respectively. Therefore, these elements are added together with B according to demand, but less than 0.005% each of the elements does not sufficiently exhibit the function.
  • Nb, Ti and V are added in the ranges of 0.005 to 0.050%, 0.005 to 0.070%, and 0.005 to 0.10%, respectively.
  • Nb has the function to fix dissolved C as NbC, and forms a nitride NbN.
  • Al and B form AlN and NB, respectively. Therefore, in order to ensure the sufficient amount of dissolved N and sufficiently decrease the amount of dissolved C, it is important to satisfy the following relations (1) and (2) : N% ⁇ 0.0015 + 14/93•Nb% + 14/27•Al% + 14/11•B% C% ⁇ 0.5•(12/93)•Nb% N/Al or N/(Al+Nb+Ti+V+B): 0.30 or more
  • Al forms AlN to decrease the amount of dissolved N.
  • N/Al In order to ensure an appropriate amount of dissolved N, N/Al must be 0.30 or more.
  • Al When Al is added in a combination with Nb, Ti, V or B, these elements also respectively form NbN, TiN, VN and BN to decrease the amount of dissolved N. Therefore, in order to ensure an appropriate amount of dissolved N, (Al+Nb+Ti+V+B) must be 0.30 or more.
  • Dissolved N 0.0010% or more
  • the content of dissolved N In order to increase the strain age hardenability of the steel sheet, the content of dissolved N must be 0.0010% or more.
  • the amount of dissolved N is determined by subtracting the amount of precipitated N from the total N amount of steel.
  • an electrolytic extraction method using a constant-potential electrolytic method is effective as the method of analyzing the amount of precipitated N.
  • an acid digestion method, a halogen method, or an electrolysis method can be used as a method of dissolving ferrite used for extraction analysis.
  • the electrolysis method can stably dissolve only ferrite without decomposing very unstable precipitates such as a carbide, a nitride, etc.
  • an acetyl-acetone system is used for electrolysis at a constant potential.
  • the results of measurement of the amount of precipitated N by constant-potential electrolysis showed best correspondence with the actual strength of parts.
  • the residue after extraction by constant-potential electrolysis is chemically analyzed to determine the amount of N in the residue.
  • the thus-determined value is considered as the amount of precipitated N.
  • the amount of dissolved N is preferably 0.0015% or more, more preferably 0.0020% or more, and most preferably 0.0030% or more.
  • the cold-rolled steel sheet of the present invention is a cold-rolled deep drawing steel sheet having excellent strain age hardenability and the above-described composition, wherein TS x r value ⁇ 750 MPa.
  • a steel sheet having a Ts x r value of less than 750 MPa cannot be widely applied to members comprising structural member components.
  • the TS x r value is preferably 850 MPa or more.
  • Conventional coating and baking conditions include 170°C for 20 min as standards.
  • a strain of 5% is applied to the steel sheet of the present invention, which contains a large amount of dissolved N, hardening can be achieved even by slow (low temperature) processing.
  • the range of aging conditions can be widened.
  • retention at a higher temperature for a longer time is advantageous as long as softening does not occurs by over aging.
  • the lower limit of the heating temperature at which hardening significantly takes place after pre-deformation is about 100°C.
  • the heating temperature of over 300°C hardening peaks, thereby causing the tendency to soften and significantly causing thermal strain and temper color.
  • the retention time is preferably 60 seconds or more.
  • retention for over 20 mines is practically disadvantageous because further hardening cannot be expected, and the production efficiency significantly deteriorates.
  • the conventional coating and baking conditions i.e., the heating temperature of 170°C and the retention time of 20 minutes
  • the heating method is not limited, and atmospheric heating with a furnace, which is generally used for coating and baking, and other methods such as induction heating, heating with a nonoxidation flame, a laser, plasma, or the like, etc. can be preferably used.
  • only a portion in which strength is desired to be increased may be selectively heated.
  • BH is 80 MPa or more
  • ⁇ TS is 40 MPa or more. More preferably, BH is 100 MPa or more, and ⁇ TS is 50 MPa or more.
  • the heating temperature in aging may be set to a higher temperature, and/or the retention time may be set to a longer time.
  • the steel sheet of the present invention has the advantage that even when the steel sheet not molded is allowed to stand at room temperature for about one year, natural aging deterioration does not occur, unlike a conventional steel sheet.
  • the cold-rolled steel sheet may be coated by hot-dip galvanization or alloying hot-dip galvanization without any problem, and TS, BH and ⁇ TS are equivalent to those before plating.
  • hot-dip galvanization and alloying hot-dip galvanization electro-galvanization, electro-tinning, electric chromium plating, electro-nickeling, and the like may be preferably used.
  • a conventional melting method such as a converter or the like
  • SRT excessively low heating temperature
  • SRT is preferably 950°C or more.
  • SRT is preferably 1300°C or less.
  • SRT is preferably 1150°C or less.
  • the total rolling reduction ratio of hot rolling starting from rough rolling to finish rolling is less than 80%, the crystal grains of the hot-rolled sheet are not sufficiently made fine. Therefore, the total rolling reduction ratio is preferably 80% or more.
  • rough rolling is preferably performed in the temperature region of the Ar 3 transformation point to 1000°C.
  • finish rolling is completed in the temperature region of over the Ar 3 transformation point, the aggregation structure is made random by ⁇ transformation to fail to obtain excellent deep drawability.
  • finish rolling is completed in the temperature region of less than the Ar 3 transformation point, a further improvement in deep drawability cannot be expected, but the rolling load is increased. Therefore, finish rolling is preferably performed in the temperature region of 600°C to the Ar 3 transformation point.
  • finish rolling is preferably performed under lubrication.
  • CT coiling temperature
  • the thus-obtained hot-rolled sheet is recrystallized and annealed by continuous annealing or batch annealing.
  • the annealing (hot-rolled sheet annealing) is carried out for recrystallizing the rolled aggregation structure formed by hot rolling in the ⁇ -phase region in finish rolling to obtain a recrystallized aggregation structure.
  • the hot-rolled sheet is cold-rolled to form a cold-rolled sheet.
  • the rolling reduction ratio of cold rolling is less than 60%, a high r value cannot be expected.
  • the rolling reduction ratio is preferably 60 to 95%.
  • the cold-rolled sheet is recrystallized and annealed.
  • the annealing is preferably carried out in either a continuous annealing line or a continuous hot-dip galvanization line.
  • the preferable annealing conditions include the annealing temperature of 650°C or more and the retention time of 5 seconds or more. When either of the annealing temperature of 650°C or more and the retention time of 5 seconds or more is not satisfied, recrystallization is not completed to deteriorate deep drawability. In order to obtain excellent deep drawability, the annealing temperature of 800°C or more and the retention time of 5 seconds or more are preferred.
  • the annealing temperature is preferably 900°C or less.
  • the cold-rolled annealed sheet obtained by recrystallizing and annealing the cold-rolled steel sheet is further coated by hot-dip galvanization or alloyed.
  • the cooling rate during the time between the completion of recrystallization annealing and the start of plating is 5°C/s or more, and the sheet temperature in hot-dip galvanization is preferably 400 to 600°C.
  • the processing temperature is preferably 400 to 600°C, and the processing time is preferably 5 to 40 seconds.
  • the cold-rolled steel sheet after recrystallization annealing or the hot-dip galvanized steel sheet may be temper-rolled for correcting the shape and controlling surface roughness.
  • the reduction ratio of temper rolling is preferably 10% or less. This is because with a rolling reduction ratio of over 10%, the r value is decreased.
  • the C content in order to control the structure to a homogeneous fine structure, and ensure a sufficient amount of an acicular ferrite phase, the C content must be 0.0015% or more. With a C content of over 0.025%, the ratio of the carbide in the steel sheet is excessively increased to significantly deteriorate ductility, the r value and moldability. Therefore, the C content is limited in the range of 0.0015 to 0.025%. From the viewpoint of improvement in moldability, the C content is preferably 0.020% or less, more preferably 0.010% or less. Particularly, from the viewpoint of stabilization of the BH amount and material properties, the C content preferably exceeds (12/93) Nb (%) (wherein Nb represents the Nb content (%)). Si: 1.0% or less
  • Si is a useful component capable of increasing the strength of the steel sheet without significantly deteriorating ductility of steel.
  • the Si content is preferably 0.10% or more.
  • Si is an element which greatly changes the transformation point during hot rolling to cause difficulties in ensuring quality and the shape, or adversely affects surface properties, chemical conversion properties, and the like, particularly the beauty of the surface of the steel sheet, and adversely affects plating properties.
  • the Si content is limited to 1.0% or less.
  • the above-described adverse effects can be kept down as long as Si is 1.0% or less.
  • Si is preferably 0.5% or less.
  • Mn 2.0% or less
  • Mn is an element effective to prevent hot cracking with S, and Mn is preferably added according to the amount of S contained. Mn also has the great effect of making fine crystal grains, and is preferably added for improving material properties. In order to stably fix S, the Mn content is preferably 0.1% or more. Mn is also an element for increasing the strength of the steel sheet, and is preferably added in an amount of 0.5% or more when higher strength is required. The Mn content is more preferably 0.8% or more.
  • the Mn content increased to this level, there is the advantage that variations in the mechanical properties of the steel sheet with respect to variations in the hot-rolling conditions, particularly strain age hardenability, are significantly improved.
  • the excessively high Mn content of over 2.0% the deformation resistance at elevated temperatures tends to increase to deteriorate weldability and weld moldability.
  • the detailed mechanism of this is not known.
  • the formation of ferrite is significantly suppressed, and the r value is significantly decreased. Therefore, the Mn content is limited to 2.0% or less. In applications required to have good corrosion resistance and moldability, the Mn content is preferably 1.5% or less. P: 0.1% or less
  • P is a useful element as a solid solution strengthening element for steel, and is preferably added in an amount of 0.002% or more from the viewpoint of an increase in strength. Particularly, when high strength is required, the P content is preferably 0.02% or more. On the other hand, when P is excessively added, steel is embrittled, and stretch-flanging properties of the steel sheet deteriorate. Also, P is liable to strongly segregate in steel, thereby causing embrittlement of a weld. Therefore, P is limited to 0.1% or less. In applications in which elongated flange processability and weld toughness are considered as important, P is preferably 0.08% or less, more preferably 0.06% or less. S: 0.02% or less
  • the Si content is as low as possible, and in the present invention, the Si content is limited to 0.02% or less.
  • S is preferably 0.015% or less.
  • S is preferably 0.010% or less.
  • Al is an element functioning as a element for deoxidation for improving cleanliness of steel, and making fine the structure of the steel sheet.
  • the Al content is preferably 0.001% or more.
  • dissolved N is used as a strengthening element, but aluminum killed steel containing Al in a suitable range has mechanical properties superior to those of conventional rimmed steel not containing Al.
  • Al is limited to 0.02% or less. From the viewpoint of stability of material quality, Al is more preferably 0.001 to 0.015%.
  • the amounts of other alloy elements are appropriately determined, and the annealing conditions are set in appropriately ranges, thereby effectively preventing coarsening.
  • N 0.0050 to 0.0250 mass %
  • N is an element for increasing the strength of the steel sheet by solid solution strengthening and strain age hardening, and in the present invention, N is the most important element.
  • an appropriate amount of N is contained, the Al content is controlled to the appropriate value, and production conditions such as the hot-rolling conditions, and the annealing conditions are controlled to ensure necessary and sufficient dissolved N in a cold-rolled product or a coated product. This exhibits the.
  • N also has the function to decrease the. transformation point, N is effective for rolling of a thin material for which rolling at a temperature greatly over the transformation point is undesirable.
  • N is limited to the range of 0.0050 to 0.0250%.
  • N is preferably in the range of 0.0070 to 0.0200%, and more preferably in the range of 0.0100 to 0.0170%.
  • the content of dissolved N (solid solution N) in the steel sheet is at least 0.0010% or more.
  • the amount of dissolved N is determined by subtracting the amount of precipitated N from the total N amount of steel.
  • electrolytic extraction analysis using constant-potential electrolysis is effective as the method of analyzing the amount of precipitated N.
  • an acid digestion method, a halogen method, or an electrolysis method can be used as the method of dissolving ferrite used for extraction analysis.
  • the electrolysis method can stably dissolve only ferrite without decomposing very unstable precipitates such as a carbide, a nitride, etc.
  • an acetyl-acetone system is used for electrolyzing at a constant potential.
  • the results of measurement of the amount of precipitated N by constant-potential electrolysis showed best correspondence with changes in actual material properties.
  • the residue after extraction by constant-potential electrolysis is chemically analyzed to determine the amount of N in the residue.
  • the thus-determined value is considered as the amount of precipitated N.
  • the amount of dissolved N is preferably 0.0020% or more. In order to obtain further high values, the amount of dissolved N is preferably 0.0030% or more. Although the upper limit of the amount of dissolved N is not limited, the mechanical properties less deteriorate even when the all amount of N remains. N/Al (the content ratio of N to Al): 0.3 or more
  • N/Al is limited to 0.3 or more.
  • N/Al is preferably 0.6 or more, and more preferably 0.8 or more.
  • Nb 0.002 to 0.050%
  • Nb effectively functions to form an acicular ferrite phase in combination with B.
  • the Nb content must be 0.002% or more.
  • a Nb content of over 0.050% the effect is saturated, and deformation resistance at elevated temperatures is significantly increased to cause difficulties in hot rolling. Therefore, Nb is limited in the range of 0.002 to 0.050%, and preferably 0.005 to 0.040%.
  • B effectively functions to form an acicular ferrite phase in combination with Nb.
  • the B content must be 0.001% or more.
  • a B content of over 0.0050% the amount of dissolved N contributing strain age hardenability is decreased. Therefore, B is limited in the range of 0.0001 to 0.0050%, and preferably 0.0003 to 0.0030%.
  • the above composition preferably further contains at least one of the following groups a to c:
  • Element of group a Cu, Ni, Cr and Mo are all contribute to an increase in strength of the steel sheet, and can be contained singly or in a combination according to demand. The effect is recognized by containing 0.01% or more each of Cu, Ni, Cr and Mo. However, with an excessively high content, deformation resistance at elevated temperatures in hot rolling is increased, or chemical conversion properties and surface treatment properties in a wide sense deteriorate, and a welded portion is hardened to deteriorate weld moldability. Therefore, Cu, Ni, Cr, and Mo are preferably contained singly at 1.0% or less, 1.0% or less, 0.5% or less, and 0.2% or less, respectively, and preferably contained in a combination at a total of 1.0% or less.
  • Ti and V are elements contributing refinement and homegenization of crystal grains, and may be added singly or in a combination according to demand. The effect can be recognized by containing 0.005% or more each of Ti and V. However, with an excessively high content, deformation resistance at elevated temperatures in hot rolling is increased, or chemical conversion properties and surface treatment properties in a wide sense deteriorate. Furthermore, there is the adverse effect of decreasing the amount of dissolved N. Therefore, Ti and V are preferably contained singly at 1.0% or less and 1.0% or less, respectively, and preferably contained in a combination at a total of 0.1% or less.
  • Elements of group c Both Ca and REM are elements useful for controlling the form of inclusions. Particularly, when the stretch flanging property is required, these elements are preferably added singly or in a combination. When the total of the elements of group d is less than 0.0010%, the effect of controlling the form of inclusions is insufficient, while when the total exceeds 0.010%, surface defects significantly occur. Therefore, the total of the elements of group d is preferably limited to the range of 0.0010 to 0.010%. This permits improvement in the stretch flanging property without causing surface defects.
  • the steel sheet of the present invention has the structure composed of an acicular ferrite phase at an area ratio of 5% or more and a ferrite phase having an average crystal grain diameter of 20 ⁇ m or less.
  • the area ratio of the acicular ferrite phase 5% or more
  • the cold-rolled steel sheet of the present invention contains the acicular ferrite phase at an area ratio of 5% or more.
  • the presence of the acicular ferrite phase at 5% or more permits the achievement of good ductility and a larger amount of strain age hardening.
  • the detailed mechanism is not known, it is thought that strain is effectively accumulated in the steel sheet during pre-strain processing before aging by the presence of the acicular ferrite phase.
  • the presence of the acicular ferrite phase improves natural aging deterioration at room temperature to be effective to obtain natural non-aging properties.
  • the area ratio of the acicular ferrite phase is preferably 10% or more.
  • the presence of a large amount of acicular ferrite phase of over 20% has the problem of deteriorating the r value. Therefore, the area ratio of the acicular ferrite phase is 5% or more, and preferably 10% to 20%.
  • the acicular ferrite phase is a low-temperature transformation phase peculiar to ultra low carbon steel having the composition of the present invention, in which no carbide is contained therein. This phase can be clearly discriminated from normal polygonal ferrite by observation on an optical microscope, and is harder than the polygonal ferrite phase because of the high internal dislocation density.
  • the acicular ferrite phase has a distribution of any one of (1) crystal grains having irregularly angular boundaries, (2) crystal grains present along crystal grains of precipitates or the like, and (3) crystal grains or crystal grain groups (many sub-boundaries are observed in relatively second phase grains) having a scratch-like pattern, singly or in a combination.
  • This acicular ferrite can be clearly distinguished from general polygonal ferrite.
  • the color tone of the corroded insides of the grains is different from martensite and bainite, and substantially the same as ordinary polygonal ferrite, and thus acicular ferrite can be clearly distinguished from martensite and bainite.
  • the acicular ferrite phase has a very high density of dislocation in the vicinities of the crystal grains and/or in the grains, and particularly the above form (3) comprises a layer having a very high dislocation density and a layer having a relatively low dislocation density.
  • the cold-rolled steel sheet of the present invention is directed to use as an automobile steel sheet required to have high moldability, and comprises a ferrite phase other than the acicular ferrite phase in order to ensure ductility.
  • the area ratio of the ferrite phase is less than 80%, it is difficult to ensure ductility and a high r value which are necessary for the automobile steel sheet required to have processability.
  • the area ratio of the ferrite phase is 80% or more, and preferably 85% or more.
  • ferrite means so-called polygonal ferrite in which no strain remains. Average crystal grain diameter of ferrite phase: 20 ⁇ m or less
  • the value used as the average crystal grain diameter is a higher one of the value calculated from a photograph of a sectional structure by a quadrature method defined by ASTM, and the nominal value determined by an intercept method defined by ASTM (refer to, for example, Umemoto et al.: Heat Treatment, 24 (1984), p334).
  • the cold-rolled steel sheet of the present invention maintains a predetermined amount of dissolved N in the product step.
  • variations in strain age hardenability occur in steel sheets containing the same amount of dissolved N, and one of the main causes of the variations is a crystal grain diameter.
  • the average crystal grain diameter is at least 20 ⁇ m or less, and preferably 15 ⁇ m or less.
  • the average crystal grain diameter of the ferrite phase is 20 ⁇ m or less, and preferably 15 ⁇ m or less.
  • the cold-rolled steel sheet of the present invention which has the above-described composition and structure, has a tensile strength (TS) of about 340 MPa to 590 MPa, a r value of as high as 1.2 or more, an excellent strain age hardenability.
  • TS tensile strength
  • the steel sheet having TS of less than 340 MPa cannot be widely applied to members each comprising a structural component.
  • TS is preferably 400 MPa or more.
  • the preferred range of the r value is 1.3 or more.
  • the conventional coating and baking conditions include 170°C and 20 min as standards.
  • a strain of 5% or more is applied to the steel sheet of the present invention containing a large amount of dissolved N, hardening can be achieved even by aging at low temperature.
  • the aging conditions can be selected from a wide range.
  • the lower limit of the heating temperature at which hardening significantly takes place after pre-deformation is about 100°C.
  • the heating temperature of over 300°C hardening peaks, thereby causing the tendency to soften and significantly causing thermal strain and temper color.
  • the retention time is preferably 60 seconds or more.
  • retention for over 20 minutes is practically disadvantageous because further hardening cannot be expected, and the production efficiency significantly deteriorates.
  • the conventional coating and baking conditions i.e., the heating temperature of 170°C and the retention time. of 20 minutes
  • the heating temperature i.e., the heating temperature of 170°C and the retention time. of 20 minutes
  • the heating method is not limited, and atmospheric heating with a furnace, which is generally used for coating and baking, and other methods such as induction heating, heating with a nonoxidation flame, a laser, plasma, or the like, etc. can be preferably used.
  • BH is 80 MPa or more (corresponding to strength in a relatively low strain region), and ⁇ TS is 40 MPa or more (corresponding to strength in a relatively high strain region). More preferably, BH is 100 MPa or more, and ⁇ TS is 50 MPa or more.
  • the heating temperature in aging may be set to a higher temperature, and/or the retention time may be set to a longer time.
  • the effect of the present invention is exhibited by a product having a relatively large thickness.
  • a sufficient cooling rate necessary for the cold-rolled sheet annealing step cannot be ensured to cause strain aging at the time of continuous annealing, thereby failing to obtain the target strain age hardenability as a product. Therefore, the steel sheet of the present invention preferably has a thickness of 3.2 mm or less.
  • the surface of the cold-rolled steel sheet may be coated by hot-dip galvanization or alloying hot-dip galvanization without any problem.
  • These coated steel sheets also exhibit TS, BH and ⁇ TS which are equivalent to those before plating.
  • the plating type any one of electro-galvanization, hot-dip galvanization, alloying hot-dip galvanization, electro-tinning, electric chromium plating, electro-nickeling, and the like may be preferably used.
  • the steel sheet of the present invention is basically produced by performing the hot rolling step in which a steel slab having the above-described composition is heated, and then roughly rolled to form a sheet bar, and the sheet bar is finish-rolled and cooled to form a coiled hot-rolled sheet, the cold rolling step in which the hot-rolled sheet is pickled and cold-rolled to form a cold-rolled sheet, and the cold-rolled sheet annealing step in which the cold-rolled sheet is continuously annealed.
  • the slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macro-segregation of components
  • an ingot making method or a thin slab casing method may be used.
  • a conventional method comprising cooling the produced slab to room temperature and then again heating the slab, or an energy-saving process of direct rolling, in which a hot slab is charged into a heating furnace without cooling and then rolled, or the slab is rolled directly immediately after being slightly kept warm, may be used without a problem.
  • direct rolling is a useful technique for effectively ensuring dissolved N.
  • the slab heating temperature is preferably 1000°C or more in order to ensure a necessary and sufficient amount of dissolved N in an initial state and satisfy the target amount of dissolved N in a product. Since a loss is increased by an increase in the oxide weight, the heating temperature is preferably 1280°C or less.
  • the slab heated under the above condition is roughly rolled to form a sheet bar.
  • the condition of the rough rolling is not defined, and rough rolling may be performed according to a normal method. However, in order to ensure the amount of dissolved N, rough rolling is preferably performed within a as short time as possible. Then, the sheet bar is finish-rolled to form a hot-rolled sheet.
  • the adjacent sheet bars are preferably bonded together during the time between rough rolling and finish rolling, and then continuously rolled.
  • a pressure welding method a laser welding method, an electron beam welding method, or the like is preferably used.
  • the ends of the coil can be stably passed, and it is thus possible to use lubricating rolling, which cannot be easily applied to ordinary single rolling for each sheet bar because of the problem of continuous rolling processes and biting property. Therefore, the rolling load can be decreased, and at the same time, the surface pressure of the roll can be decreased, thereby increasing the life time of the roll.
  • the temperature distributions of the sheet bar in the width direction and the long direction thereof are preferably made uniform by using any one or both of a sheet bar edge heater for heating the ends of the sheet bar in the width direction and a sheet bar heater for heating the ends of the sheet bar in the long direction. This can further decrease the variations in material properties of the steel sheet.
  • the sheet bar edge heater and the sheet bar heater are preferably of an induction heating type.
  • the difference in temperature in the width direction is first corrected by the sheet bar edge heater.
  • the heating amount depends upon the steel composition, the temperature distribution in the width direction at the finisher entrance is preferably set in the range of about 20°C or less.
  • the difference in temperature in the long direction is corrected by the sheet bar heater.
  • the heating amount is preferably set so that the temperatures at the ends in the long direction are about 20°C higher than the temperature at the center. Finisher delivery temperature: 800°C or more
  • the finisher deliver temperature FDT is 800°C or more.
  • the structure of the steel sheet becomes inhomogeneous, and the processed structure partially remains to leave heterogeneity of the structure after the cold-rolled sheet annealing step. Therefore, the danger of producing various troubles in press forming is increased.
  • a high coiling temperature is used for avoiding the processed micro structure from remaining, coarse crystal grains are produced to cause the same troubles as described above. With the high coiling temperature, the amount of dissolved N is significantly decreased to cause difficulties in obtaining a target tensile strength of 340 MPa or more. Therefore, the finisher deliver temperature FDT is 800°C or more.
  • the FDT is preferably 820°C or more. In order to improve the r value, the FDT is preferably the Ac 3 transformation point or more. Although the upper limit of FDT is not limited, a scale scar significantly occurs at excessively high FDT.
  • the FDT is preferably up to about 1000°C. Coiling temperature: 800°C or less
  • the strength of the steel sheet is liable to increase as the coiling temperature CT decreases.
  • the CT is preferably 800°C or less. With a CT of less than 200°C, the shape of the steel sheet is readily disturbed to increase the. danger of causing troubles in a practical operation, thereby deteriorating homogeneity of material properties. Therefore, the CT is preferably 200°C or more.
  • the CT is preferably 300°C or more, and more preferably 350°C or more.
  • lubricating rolling in finish rolling, may be performed for decreasing the hot rolling load.
  • the lubricating rolling has the effect of further making homogeneous the shape and material properties of the hot-rolled sheet.
  • the frictional coefficient is preferably in the range of 0.25 to 0.10.
  • the hot-rolled sheet subjected to the above-described hot rolling step is then pickled and cold-rolled in the cold rolling step to form a cold-rolled sheet.
  • the pickling conditions may be the same as conventional known conditions, and are not limited. When the scale of the hot-rolled sheet is extremely small, cold rolling may be immediately after hot rolling without pickling.
  • the cold rolling conditions may be the same as conventional known conditions, and are not limited.
  • the reduction ratio of cold rolling is preferably 60% or more. The reasons for limiting the conditions of the cold-rolled sheet annealing step are described below.
  • the cold-rolled sheet is then subjected to the cold-rolled sheet annealing step comprising continuous annealing and cooling.
  • Continuous annealing temperature temperature in the ferrite-austenite two-phase coexistence region
  • the acicular ferrite phase is formed.
  • the (111) aggregation structure is strongly developed in the ferrite phase to obtain a high r value.
  • the annealing temperature of continuous annealing is limited to the recrystallization temperature or more in the ferrite-austenite two-phase coexistence region.
  • the temperature is preferably set so that the fraction of austenite is 10% to 50%.
  • the annealing temperature is preferably the recrystallization temperature or more.
  • the retention time of continuous annealing is preferably as short as possible in order to ensure the production efficiency, the fine structure and the amount of dissolved N. From the viewpoint of stability of the operation, the retention time is preferably 10 seconds or more. Also, in order to ensure the fine structure and the amount of dissolved N, the retention time is preferably 90 seconds or less. From the viewpoint of stability of material properties, the retention time is preferably 20 seconds or more. Cooling after continuous annealing: cooling to the temperature region of 500°C or less at a cooling rate of 10 to 300°C/s
  • Cooling after soaking by continuous annealing is important for making fine the structure, forming the acicular ferrite phase, and ensuring the amount of dissolved N.
  • cooling is continuously carried out to the temperature region of at least 500°C or less at a cooling rate of 10°C/s or more.
  • a cooling rate of less than 10°C/s a necessary amount of acicular ferrite phase, a homogeneous fine structure and a sufficient amount of dissolved N cannot be obtained.
  • a cooling rate of over 300°C/s homogeneity in material properties of the steel sheet in the width direction is insufficient.
  • the stop temperature of cooling at a cooling rate of 10 to 300°C/s after continuous annealing exceeds 500°C, refinement of the structure cannot be attained.
  • Temper rolling or lever processing elongation of 0.5 to 10%
  • temper rolling or leveler processing may be carried out subsequent to the cold rolling step.
  • the total elongation of temper rolling or leveler processing is less than 0.5%, the desired purpose of correcting the shape and controlling roughness cannot be achieved.
  • the elongation is preferably 5% or less. It is confirmed that the processing system of temper rolling is different from that of leveler processing, but the effects of both processes are substantially the same. Temper rolling and leveler processing are effective even after plating.
  • C is an element for increasing the strength of a steel sheet, and 0.025% or more of C must be contained for controlling the structure to a homogeneous fine structure, which is an important requirement of the present invention, and ensuring a sufficient amount of a martensite phase.
  • a C content of over 0.15% the ratio of the carbide in the steel sheet is excessively increased to significantly deteriorate ductility and moldability.
  • the C content is limited in the range of 0.025 to 0.15%.
  • the C content is preferably 0.08% or less.
  • the C content is preferably 0.05% or less.
  • Si is a useful component capable of increasing the strength of the steel sheet without significantly deteriorating ductility of steel.
  • the Si content is preferably 0.005% or more, and more preferably 0.10% or more.
  • Si is an element which greatly changes the transformation point during hot rolling to cause difficulties in ensuring quality and the shape, or adversely affects surface properties, chemical conversion properties, and the like, particularly the beauty of the surface of the steel sheet, and adversely affects plating properties.
  • the Si content is limited to 1.0% or less.
  • the above-described adverse effects can be kept down as long as Si is 1.0% or less.
  • Si is preferably 0.5% or less.
  • Mn 2.0% or less
  • Mn is an element effective to prevent hot cracking with S, and Mn is preferably added according to the amount of S contained. Mn also has the great effect of making fine crystal grains, and is preferably added for improving material properties. Furthermore, Mn is an element effective to stably form martensite during rapid cooling after continuous annealing. In order to stably fix S, the Mn content is preferably 0.2% or more. Mn is also an element for increasing the strength of the steel sheet, and is preferably added in an amount of 1.2% or more when a strength TS of over 500 MPa is required. The Mn content is more preferably 1.5% or more.
  • the Mn content is increased to this level, there is the advantage that variations in the mechanical properties of the steel sheet with respect to variations in the hot-rolling conditions, particularly strain age hardenability, are significantly improved.
  • the Mn content is limited to 2.0% or less.
  • the Mn content is preferably 1.7% or less. P: 0.08% or less
  • P is a useful element as a solid solution strengthening element for steel, and is preferably added in an amount of 0.001% or more, and more preferably 0.015% or more, from the viewpoint of an increase in strength.
  • P is excessively added, steel is embrittled, and stretch-flanging properties of the steel sheet deteriorate. Also, P is liable to strongly segregate in steel, thereby causing embrittlement of a weld. Therefore, P is limited to 0.08% or less. In applications in which elongated flange processability and weld toughness are considered as important, P is preferably 0.04% or less. S: 0.02% or less
  • the Si content is as low as possible, and in the present invention, the S content is limited to 0.02% or less.
  • S is preferably 0.015% or less.
  • S is preferably 0.008% or less.
  • Al is an element functioning as a deoxidization for improving cleanliness of steel, and making fine the structure of the steel sheet.
  • the Al content is preferably 0.001% or more.
  • dissolved N is used as a strengthening element, but aluminum killed steel containing Al in a suitable range has mechanical properties superior to those of conventional rimmed steel not containing Al.
  • Al is limited to 0.02% or less. From the viewpoint of stability of material properties, Al is more preferably 0.001 to 0.015%.
  • N is an element for increasing the strength of the steel sheet by solid solution strengthening and strain age hardening, and in the present invention, N is the most important element.
  • an appropriate amount of N is contained, the Al content is controlled to the appropriate value, and production conditions such as the hot-rolling conditions, and the annealing conditions are controlled to ensure necessary and sufficient dissolved N in a cold-rolled product or a coated product.
  • N also has the function to decrease the transformation point, N is effective for rolling of a thin material for which rolling at a temperature greatly over the transformation point is undesirable.
  • N is limited to the range of 0.0050 to 0.0250%. From the viewpoint of improvement in stability of material properties and yield over the entire production process, N is preferably in the range of 0.0070 to 0.0170%. With the N amount in the range of the present invention, there is no adverse effect on weldability, and the like. Dissolved N: 0.0010% or more
  • the content of dissolved N (solid solution N) in the steel sheet is at least 0.0010% or more.
  • the amount of dissolved N is determined by subtracting the amount of precipitated N from the total N amount of steel.
  • electrolytic extraction analysis using constant-potential electrolysis is effective as the method of analyzing the amount of. precipitated N.
  • an acid digestion method, a halogen method, or an electrolysis method can be used as the method of dissolving ferrite used for extraction analysis.
  • the electrolysis method can stably dissolve only ferrite without decomposing very unstable precipitates such as a carbide, a nitride, etc.
  • an acetyl-acetone system is used for electrolysis at a constant potential.
  • the results of measurement of the amount of precipitated N by constant-potential electrolysis showed best correspondence with changes in actual material properties.
  • the residue after extraction by constant-potential electrolysis is chemically analyzed to determine the amount of N in the residue.
  • the thus-determined value is considered as the amount of precipitated N.
  • the amount of dissolved N is preferably 0.0020% or more, more preferably 0.0020% or more. In order to obtain further high values, the amount of dissolved N is preferably 0.0030% or more. Although the upper limit of the amount of dissolved N is not limited, the mechanical properties less deteriorate even when the all amount of N added remains. N/Al (the content ratio of N to Al): 0.3 or more
  • the above component preferably further contains at least one of the following groups d to g:
  • Elements of group d are all contribute to an increase in strength of the steel sheet, and can be contained singly or in a combination according to demand. The effect is recognized by containing 0.005% or more each of Cu, Ni, Cr and Mo.
  • the elements of group a are preferably contained in a total of 1.0% or less. With a Mo content of 0.05% or more, the r value is significantly decreased in some cases. In the present invention, therefore, the Mo content is preferably limited to less than 0.05%.
  • Nb, Ti and V are elements contributing refinement and homegenization of crystal grains, and may be added singly or in a combination according to demand. The effect can be recognized by containing 0.005% or more each of Nb, Ti and V. However, with an excessively high content, deformation resistance in hot rolling at elevated temperatures is increased, or chemical conversion properties and surface treatment properties in a wide sense deteriorate. Therefore, the elements in group b are preferably contained at a total of 0.1% or less.
  • B is an element having the effect of improving hardenability of steel, and can be contained for increasing the fraction of a low-temperature transformation phase other than the ferrite phase to increase strength of steel according to demand. This effect is recognized with a B content of 0.0005% or more. However, with an excessively high B content, deformability at elevated temperatures in hot rolling deteriorates to produce BN, decreasing the amount of dissolved N. Therefore, the B content is preferably 0.0030% or less.
  • Both Ca and REM are elements useful for controlling the form of inclusions. Particularly, when the stretch flanging property is required, these elements are preferably added singly or in a combination. When the total of the elements of group d is less than 0.0010%, the effect of controlling the form of inclusions is insufficient, while when the total exceeds 0.010%, surface defects significantly occur. Therefore, the total of the elements of group d is preferably limited to the range of 0.0010 to 0.010%. This permits improvement in the stretch flanging property without causing surface defects.
  • the cold-rolled steel sheet of the present invention is directed to use as an automobile steel sheet required to have some extent of moldability, and has a structure containing the ferrite phase at an area ratio of 80% or more in order to ensure ductility.
  • the ferrite phase With the ferrite phase at an area ratio of less than 80%, it is difficult to ensure ductility required for an automobile steel sheet required to have moldability.
  • the area ratio of the ferrite phase is preferably 85% or more.
  • "ferrite” means so-called polygonal ferrite in which no strain remains. Average crystal grain diameter of ferrite phase: 10 ⁇ m or less
  • the value used as the average crystal grain diameter is a higher value of the value calculated from a photograph of a sectional structure by a quadrature method defined by ASTM, and the nominal value determined by an intercept method defined by ASTM (refer to, for example, Umemoto et al.: Heat Treatment, 24 (1984), p334).
  • the cold-rolled steel sheet of the present invention maintains a predetermined amount of dissolved N in the product step.
  • variations in strain age hardenability occur in steel sheets containing the same amount of dissolved N, and one of the main causes of the variations is a crystal grain diameter.
  • the average crystal grain diameter is at least 10 ⁇ m or less, and preferably 8 ⁇ m or less.
  • the average crystal grain diameter of the ferrite phase is 10 ⁇ m or less, and preferably 8 ⁇ m or less.
  • the structure of the present invention contains the ferrite phase with an average crystal grain diameter of 10 ⁇ m or less at an area ratio of 80% or more.
  • Area ratio of martensite phase 2% or more
  • the cold-rolled steel sheet of the present invention contains the martensite phase as a second phase at an area ratio of 2% or more.
  • the presence of 2% or more of the martensite phase can produce good ductility and a large amount of strain age hardening.
  • this effect is supposed to be due to the effective accumulation of strain in the steel sheet due to the presence of the martensite phase during pre-strain processing before aging.
  • the presence of the martensite phase is effective to improve aging deterioration.
  • the area ratio of the martensite phase is preferably 5% or more.
  • the presence of the martensite phase at an area ratio of over 20% causes the problem of deteriorating ductility. Therefore, the area ratio of the martensite phase is 2% or more, and preferably 5% to 20%.
  • the cold-rolled steel sheet of the present invention which has the above-described composition and structure, has a tensile strength (TS) of 440 MPa to about 780 MPa, a high r value of 1.2 or more obtained by controlling the aggregation structure of the ferrite base phase, and excellent strain age hardenability.
  • TS tensile strength
  • a steel sheet having TS of less than 440 MPa cannot be widely applied to members having structural components.
  • TS is preferably 500 MPa or more.
  • the preferable range of the r value is 1.4 or more.
  • the amount of pre-strain is an important factor.
  • the deformation stress in the above-described deformation system can be referred to as an amount of approximately uniaxial strain (tensile strain) except the case of excessive deep drawing, (2) the amount of uniaxial strain of an actual part exceeds 5%, and (3) the strength of a part sufficiently corresponds to the strength (YS and TS) obtained after strain aging with a pre-strain of 5%.
  • the pre-deformation of strain aging is defined to a tensile strain of 5%.
  • Conventional coating and baking conditions include 170°C and 20 min as standards.
  • a strain of 5% is applied to the steel sheet of the present invention, which contains a large amount of dissolved N, hardening can be achieved even by aging at low temperature. In other words, the range of aging conditions can be widened.
  • retention at a higher temperature for a longer time is advantageous as long as softening does not occurs by over aging.
  • the lower limit of the heating temperature at which hardening significantly takes place after pre-deformation is about 100°C.
  • the heating temperature of over 300°C hardening peaks, thereby causing the tendency to soften and significantly causing thermal strain and temper color.
  • the retention time is preferably 60 seconds or more.
  • retention for over 20 mines is practically disadvantageous because further hardening cannot be expected, and the production efficiency significantly deteriorates.
  • the conventional coating and baking conditions i.e., the heating temperature of 170°C and the retention time of 20 minutes
  • the heating temperature i.e., the heating temperature of 170°C and the retention time of 20 minutes
  • the heating method is not limited, and atmospheric heating with a furnace, which is generally used for coating and baking, and other methods such as induction heating, heating with a nonoxidation flame, a laser, plasma, or the like, etc. can be preferably used.
  • BH is 80 MPa or more
  • ⁇ TS is 40 MPa or more. More preferably, BH is 100 MPa or more, and ⁇ TS is 50 MPa or more.
  • the heating temperature in aging may be set to a higher temperature, and/or the retention time may be set to a longer time.
  • the steel sheet of the present invention has the advantage that when the steel sheet is allowed to stand at room temperature for about one week without heating after forming, an increase in strength of about 40% of that at the time of complete aging can be expected.
  • the steel sheet of the present invention also has the advantage that even when it is allowed in an unmolded state at room temperature for a long time, aging deterioration (an increase in YS and a decrease in El (elongation)) does not occurs, unlike a conventional aging steel sheet.
  • aging deterioration an increase in YS and a decrease in El (elongation)
  • an increase in YS is 30 MPa or less
  • a decrease in elongation is 2% or less
  • a recovery of yield point elongation is 0.2% or less.
  • the surface of the cold-rolled steel sheet may be coated by hot-dip galvanization or alloying hot-dip galvanization without any problem, and TS, BH and ⁇ TS are equivalent to those before plating.
  • TS, BH and ⁇ TS are equivalent to those before plating.
  • electro-galvanization, hot-dip galvanization, alloying hot-dip galvanization, electro-tinning, electric chromium plating, electro-nickeling, and the like may be preferably used.
  • the steel sheet of the present invention is basically produced by performing the hot rolling step in which a steel slab having the above-described composition is heated, and then roughly rolled to form a sheet bar, and the sheet bar is finish-rolled and cooled to form a coiled hot-rolled sheet, the cold rolling step in which the hot-rolled sheet is pickled and cold-rolled to form a cold-rolled sheet, and the cold-rolled sheet annealing step in which the cold-rolled sheet is box-annealing and then continuously annealed.
  • the slab used in the production method of the present invention is preferably produced by a continuous casting method in order to prevent macro-segregation of components
  • an ingot making method or a thin slab casing method may be used.
  • a conventional method comprising cooling the produced slab to room temperature and then again heating the slab, or an energy-saving process of direct rolling comprising charging a hot slab into a heating furnace without cooling and then rolling it, or rolling directly the slab immediately after slightly keeping it warm may be used without a problem.
  • direct rolling is a useful technique for effectively ensuring dissolved N.
  • the slab heating temperature is preferably 1000°C or more in order to ensure a necessary and sufficient amount of dissolved N in an initial state and satisfy the target amount of dissolved N in a product. Since a loss is increased by an increase in the oxide weight, the heating temperature is preferably 1280°C or less.
  • the slab heated under the above condition is roughly rolled to form a sheet bar.
  • the condition of the rough rolling is not defined, and rough rolling may be performed according to a normal method. However, in order to ensure the amount of dissolved N, rough rolling is preferably performed within a as short time as possible. Then, the sheet bar is finish-rolled to form a hot-rolled sheet.
  • the adjacent sheet bars are preferably bonded together during the time between rough rolling and finish rolling, and then continuously rolled.
  • a pressure welding method a laser welding method, an electron beam welding method, or the like is preferably used.
  • the ends of the coil can be stably passed, and it is thus possible to use lubricating rolling, which cannot be easily applied to ordinary single rolling for each sheet bar because of the problem of continuous rolling processes and biting property. Therefore, the rolling load can be decreased, and at the same time, the surface pressure of the roll can be decreased, thereby increasing the life time of the roll.
  • the temperature distributions of the sheet bar in the width direction and the long direction thereof are preferably made uniform by using any one or both of a sheet bar edge heater for heating the ends of the sheet bar in the width direction and a sheet bar heater for heating the ends of the sheet bar in the long direction. This can further decrease the variations in material properties of the steel sheet.
  • the sheet bar edge heater and the sheet bar heater are preferably of an induction heating type.
  • the difference in temperature in the width direction is first corrected by the sheet bar edge heater.
  • the heating amount depends upon the steel composition, the temperature distribution in the width direction at the finisher entrance is preferably set in the range of about 20°C or less.
  • the difference in temperature in the long direction is corrected by the sheet bar heater.
  • the heating amount is preferably set so that the temperatures at the ends in the long direction are about 20°C higher than the temperature at the center. Finisher delivery temperature: 800°C or more
  • the finisher deliver temperature FDT is 800°C or more.
  • the structure of the steel sheet becomes inhomogeneous, and the processed structure partially remains to leave heterogeneity of the structure after the cold-rolled sheet annealing step. Therefore, the danger of causing various troubles in press forming is increased.
  • the finisher deliver temperature FDT is 800°C or more.
  • the FDT is preferably 820°C or more.
  • the upper limit of FDT is not limited, a scale scar significantly occurs at excessively high FDT.
  • the FDT is preferably up to about 1000°C.
  • cooling after finish rolling is not strictly limited, the conditions described below are preferable from the viewpoint of homogeneity in material properties of the steel sheet in the long direction and the width direction thereof.
  • cooling is preferably started immediately after (within 0.5 seconds after) finish rolling, and the mean cooling rate in cooling is preferably 40°C/s or more.
  • the steel sheet can be rapidly cooled in the high temperature region where AlN precipitates to effectively ensure N in a solid solution state.
  • the starting time of cooling or the cooling rate does not satisfy the above condition, grain growth excessively proceeds to fail to achieve fine crystal grains, and promote AlN precipitation due to stain energy introduced by rolling. Therefore, the amount of dissolved N tends to decrease, and the structure tends to be made inhomogeneous.
  • the cooling rate is preferably kept at 300°C/s or less. Coiling temperature: 800°C or less
  • the strength of the steel sheet is liable to increase as the coiling temperature CT decreases.
  • the CT is preferably 800°C or less. With a CT of less than 200°C, the shape of the steel sheet is readily disturbed to increase the danger of causing troubles in a practical operation, thereby deteriorating homogeneity of material properties. Therefore, the CT is preferably 200°C or more.
  • the CT is preferably 300°C or more, and more preferably 350°C or more.
  • lubricating rolling may be performed for decreasing the hot rolling load.
  • the lubricating rolling has the effect of further making homogeneous the shape and material properties of the hot-rolled sheet.
  • the frictional coefficient is preferably in the range of 0.25 to 0.10.
  • the hot-rolled sheet subjected to the above-described hot rolling step is then pickled and cold-rolled in the cold rolling step to form a cold-rolled sheet.
  • the pickling conditions may be the same as conventional known conditions, and are not limited. When the scale of the hot-rolled sheet is extremely small, cold rolling may be immediately after hot rolling without pickling.
  • the cold rolling conditions may be the same as conventional known conditions, and are not limited.
  • the reduction ratio of cold rolling is preferably 40% or more. The reasons for limiting the conditions of the cold rolling step are described below.
  • the cold-rolled sheet is then subjected to the cold-rolled sheet annealing step comprising box annealing and continuous annealing.
  • Box annealing temperature the recrystallization temperature to 800°C
  • the cold-rolled sheet is subjected to box annealing to control the aggregation structure of the ferrite phase as a base.
  • the r value of the produced sheet can be increased.
  • box annealing the (111) aggregation structure suitable for increasing the r value is readily formed in the produced sheet.
  • box annealing is preferably performed in an annealing atmosphere containing a nitrogen gas as a main component and 3 to 5% of hydrogen gas.
  • the heating and cooling rates may be the same as normal box annealing, and are about 30°C/hr. By using 100% hydrogen gas as an annealing atmosphere gas, the higher heating and cooling rates may be used.
  • Continuous annealing temperature Ac 1 transformation point to (Ac 3 transformation point - 20°C)
  • the continuous annealing temperature is preferably Ac 1 transformation point to (Ac 3 transformation point - 20°C).
  • the retention time of continuous annealing is preferably as short as possible in order to ensure the production efficiency, the fine structure and the amount of dissolved N. -From the viewpoint of stability of the operation, the retention time is preferably 10 seconds or more.
  • the retention time is preferably 120 seconds or less. From the viewpoint of stability of material properties, the retention time is preferably 20 seconds or more. Cooling after continuous annealing: cooling to the temperature region of 500°C or less at a cooling rate of 10 to 300°C/s
  • Cooling after soaking by continuous annealing is important for making fine the structure, forming the martensite phase, and ensuring the amount of dissolved N.
  • cooling is continuously carried out to the temperature region of at least 500°C or less at a cooling rate of 10°C/s or more.
  • a cooling rate of less than 10°C/s a necessary amount of martensite phase, a homogeneous fine structure and a sufficient amount of dissolved N cannot be obtained.
  • a cooling rate of over 300°C/s homogeneity in material properties of the steel sheet in the width direction deteriorates due to a significant increase in the amount of supersaturated dissolved C.
  • Over aging condition retention in the temperature region of 350°C to the cooling stop temperature for 20 seconds or more subsequent to cooling after continuous annealing
  • Over aging may be performed by retention in the temperature region of 350°C to the cooling stop temperature for 20 seconds or more subsequent to the stop of cooling after soaking by continuous annealing.
  • the amount of dissolved C can be selectively decreased, while the amount of dissolved N is maintained.
  • the temperature region is preferably 350°C or more.
  • the retention time is preferably 120 second or less.
  • continuous annealing after box annealing can be performed in a continuous hot-dip coating line comprising hot-dip galvanization subsequent to cooling after continuous annealing or further alloying to produce a hot-dip galvanized steel sheet.
  • Temper rolling or lever processing elongation of 0.2 to 15%
  • temper rolling or leveler processing may be carried out subsequent to the cold rolling step.
  • the total elongation of temper rolling or leveler processing is less than 0.2%, the desired purpose of correcting the shape and controlling roughness cannot be achieved.
  • ductility significantly deteriorates. It is confirmed that the processing system of temper rolling is different from that of leveler processing, but the effects of both processes are substantially the same. Temper rolling and leveler processing are effective after plating.
  • the difference ( ⁇ TS) between the tensile strength of the specimen after application of tensile strain and heat treatment and the tensile strength of a product is defined as the strength increasing ability of heat treatment.
  • the amount of strain introduced by forming, or the heat treatment temperature after processing is preferably as high as possible.
  • the strength can be sufficiently increased even by heat treatment at a temperature lower than conventional heat treatment, i.e., a temperature of 200°C or less, after forming.
  • a heat treatment temperature of less than 120°C the strength cannot be sufficiently increased with the low train applied.
  • the heat treatment temperature of over 350°C after forming softening proceeds. Therefore, the temperature of heat treatment after forming is preferably about 120 to 350°C.
  • the heating method is not limited, and hot gas heating, infrared furnace heating, hot-bath heating, direct current heating, induction heating, and the like can be used. Alternatively, only a portion where strength is desired to be increased is selectively heated.
  • the amount of dissolved N, the microstructure, tensile properties, the r value, strain age hardenability, and aging property were examined.
  • the examination methods were as follows:
  • the amount of dissolved N was determined by subtracting the amount of precipitated N from the total N amount of steel determined by chemical analysis.
  • the amount of precipitated N was determined by an analysis method using a constant-potential electrolytic method.
  • a test specimen was obtained from each of cold-rolled annealed steel sheets, and the microstructure of a section (C section) perpendicular to the rolling direction was imaged with an optical microscope or a scanning electron microscope. Then, the fraction of the ferrite texture and the type and the structure fraction of a second phase were determined by an image analysis apparatus.
  • the value used as the average crystal grain diameter was a higher one of the value calculated from a photograph of a sectional structure by a quadrature method defined by ASTM, and the nominal value determined from a photograph of a sectional structure by an intercept method defined by ASTM (refer to, for example, Umemoto et al.: Heat Treatment, 24 (1984), p334).
  • test specimen of JIS No. 5 was obtained from each of cold-rolled annealed steel sheets in the rolling direction, and a tensile test was carried out with a strain rate of 3 x 10 -3 /s according to the regulations of JIS Z 2241 to determine yield stress YS, tensile strength TS, and elongation El.
  • YS5% represents deformation stress in 5% pre-deformation of the produced sheet
  • YSBH and TSBH represent yield stress and tensile strength, respectively, after pre-deformation and heat treatment
  • TS represents the tensile strength of the produced sheet.
  • a test specimen of JIS No. 5 was obtained from each of the cold-rolled annealed steel sheets in each of the rolling direction (L direction), the direction (D direction) at 45° with the rolling direction, and the direction (C direction) at 90° with the rolling direction.
  • rmean (rL + 2rD + rD)/4 wherein rL represents the r value in the rolling direction (L direction), rD represents the r value in the direction (D direction) at 45° with the rolling direction, and rL represents the r value in the direction (C direction) at 90° with the rolling direction.
  • test specimen of JIS No. 5 was obtained from each of cold-rolled annealed steel sheets in the rolling direction, and then subjected to aging at 50°C for 200 hours, followed by a tensile test.
  • the difference in yield elongation ⁇ Y-E1 between before and after aging was determined from the obtained results to evaluate aging properties at normal temperature. When ⁇ Y-El was zero, it was evaluated that the specimen has non-aging properties and excellent natural aging resistance.
  • test specimen of JIS No. 5 was obtained from each of produced sheets in the rolling direction, and then a pre-strain of 10% was applied thereto. Then, heat treatment was conducted for 20 minutes at a conventional heat treatment temperature of 120°C and a temperature of 170°C corresponding to coating and baking, and then tensile strength was determined.
  • the decrease ( ⁇ E1) in total elongation by natural aging was determined as the difference between the total elongation measured with a specimen of JIS N0 5 obtained from the produced sheet in the rolling direction, and the total elongation measured with a specimen of JIS N0 5 separately obtained from the produced sheet in the rolling direction after accelerated aging (retention at 100°C for 8 hours) of natural aging.
  • a steel slab having each of the compositions shown in Table 1 was hot-rolled sheet having a thickness of 3.5 mm, and then cold-rolled to a cold-rolled sheet having a thickness of 0.7 mm under the conditions shown in Table 2. Then, the cold-rolled sheet was recrystallized, annealed and further galvannealed in a continuous annealing line or a continuous annealing and galvanizing line. Then, the annealed sheet was temper-rolled with a rolling reduction ratio of 1.0 % to produce a cold-rolled steel sheet and a galvannealed steel sheet having both sides coated with a weight of 45 g/m 2 per side. In Table 2, the finisher deliver temperatures of the others are the Ar 3 transformation point or more.
  • Table 3 indicates that with all the cold-rolled steel sheets and the galvannealed steel sheets obtained according to the present invention, a high r value and excellent strain age hardenability are obtained, as compared with comparative examples. Particularly, in the suitable examples in which the crystal grain diameter is 20 ⁇ m or less, the decrease in elongation due to natural aging is also as low as 2.0% or less.
  • a slab of steel symbol B shown in Table 1 was hot-rolled under the same production conditions as No. 2 shown in Table 2 in which the heating temperature was 1100°C, and the finisher deliver temperature of hot rolling was 900°C, and then coiled at coiling temperature of 550°C into a coil.
  • the thus-obtained coil was cold-rolled with a reduction ratio of 80%, and then recrystallized and annealed at 840°C.
  • tensile strength TS was 365 MPa, and the r value was 1.7.
  • a test specimen of JIS No. 5 was obtained from the cold-rolled steel sheet in the rolling direction, and a tensile strain of 10% was applied by a tensile test machine.
  • Table 4 indicates that the increase in strength increases as the heat treatment temperature increases, and the heat treatment time increases.
  • a sufficient increase in tensile strength of 82 MPa (85% or more of an increase in heat treatment for 20 minutes) can be obtained even by heat treatment at low temperature of 120°C for a short retention time of 2 minutes. It is thus found that with the steel sheet of the present invention, good strain age hardenability can be obtained even by heat treatment at a low temperature for a short time.
  • heat treatment at a normal temperature for a normal time causes no problem. It was confirmed that with the galvanized steel sheets and the galvannealed steel sheets obtained by hot-dip galvanizing and heat alloying the cold-rolled sheets, the same results as shown in Table 4 are obtained.
  • a steel slab having each of the compositions shown in Table 6 was hot-rolled under the conditions shown Table 7 to form a hot-rolled sheet having a thickness of 3.5 mm.
  • Each of the thus-obtained hot-rolled sheets was cold-rolled under the conditions shown in Table 7 to form a cold-rolled sheet having a thickness of 0.7 mm, and then recrystallized and annealed under the conditions shown in the same table.
  • Some of the annealed sheets were further coated by hot-dip galvanization or alloying hot-dip galvanization under the conditions shown in the same table.
  • the thus-obtained produced sheets were examined with respect to the amount of dissolved N, the microstructure, tensile properties, and strain age hardenability.
  • Table 8 indicates that with the steel sheets of the present invention, TS x r value ⁇ 750 MPa (in a combination with at least one of B, Nb, Ti and V, Ts x r value 850 ⁇ MPa), BH ⁇ 80 MPa and ⁇ TS ⁇ 40 MPa are satisfied, while in the comparative examples, at least one of the three properties does not reach the level of the present invention.
  • each of the thus-obtained hot-rolled sheets was cold-rolled by the cold rolling step under the conditions shown in Table 10 to form a cold-rolled sheet. Then, each of the cold-rolled sheets was continuously annealed under the conditions shown in Table 10. Some of the cold-rolled sheets were further temper-rolled after the cold-rolled sheet annealing step.
  • the thus-obtained cold-rolled annealed sheets were examined with respect to the amount of dissolved N, the microstructure, tensile properties, the r value, strain age hardenability and aging properties.
  • each of the steel sheets of Nos. 4 and 10 was hot-dip galvanized to form a coated steel sheet, and evaluated with respect to the same properties as described above.
  • All examples of the present invention have excellent ductility, an extremely high BH amount and ⁇ TS, excellent strain age hardenability, a mean r value of as high as 1.2 or more, and non-aging properties at natural aging (excellent natural aging resistance).
  • the mean r value is decreased by 0.2, and the elongation El is decreased by about 1%, as compared with the cold-rolled steel sheets, because of shrinkage restriction of the coated layer in the width direction.
  • strain age hardenability and natural aging resistance are substantially the same as those before coating.
  • the comparative examples out of the range of the present invention ductility deteriorates, the BH amount and ⁇ TS are low, or natural aging deterioration significantly occurs. Therefore, the comparative examples do not have all the intended properties, and thus cannot be said as steel sheets having sufficient properties.
  • Steel sheet No. 11 contains C, Al, N, and N/Al out of the range of the present invention, and thus the r value, the BH amount, ⁇ TS and natural aging resistance deteriorate.
  • Steel sheet No. 12 contains B and Nb out of the range of the present invention, and thus the amount of acicular ferrite is greatly deviated from the range of the present invention, deteriorating the BH amount, ⁇ TS, and natural aging resistance.
  • Steel sheet No. 13 contains B out of the range of the present invention, and thus the amount of acicular ferrite is greatly deviated from the range of the present invention, deteriorating the r value, the BH amount, ⁇ TS, and natural aging resistance.
  • Steel sheet No. 14 contains Nb out of the range of the present invention, and thus the amount of dissolved N is greatly lower than the range of the present invention, deteriorating strain age hardenability.
  • Steel sheet No. 15 contains N out of the range of the present invention, and thus the amount of dissolved N is low, deteriorating strain age hardenability.
  • the hot rolling conditions and the cold-rolled sheet annealing conditions are deviated from the suitable ranges, and thus the microstructure is out of the range of the present invention, decreasing the BH amount and ⁇ TS and deteriorating strain age hardenability and natural aging resistance.
  • Example 12 Steel having the composition shown in Table 12 was formed in a slab by the same method as Example 4, and then heated and temper-rolled under the conditions shown in Table 13 to form a sheet bar.
  • the sheet bar was then hot-rolled by the hot rolling step comprising finish rolling under the conditions shown in Table 13 to form a hot-rolled sheet.
  • the adjacent sheet bars on the finisher entrance side after rough rolling were bonded together by a melt welding method and then continuously rolled.
  • the temperatures of the ends of the sheet bar were controlled in the width direction and the length direction by using an induction heating-type sheet bar edge heater and a sheet bar heater.
  • the thus-obtained hot-rolled sheet was cold-rolled by the cold rolling step comprising pickling and cold rolling under the conditions shown in Table 13 to form a cold-rolled sheet having a thickness of 1.6 mm. Then, the cold-rolled sheet was continuously annealed under the conditions shown in Table 13.
  • the thus-obtained cold-rolled annealed sheet was examined with respect to the amount of dissolved N, the microstructure, tensile properties, the r value, and strain age hardenability by the same methods as Example 4.
  • the tensile property of each cold-rolled annealed sheet was measured at ten positions in each of the width direction and the long direction to examine variations in yield strength, tensile strength and elongation.
  • All the examples of the present invention have excellent strain age hardenability and a high r value, and exhibit extremely high stable BH amount, ⁇ TS and mean value regardless of variations in production conditions. It was also recognized that in the examples of the present invention, by performing continuous rolling and controlling the temperature of the sheet bar in the long direction and the width direction, the thickness precision and the shape of the produced steel sheet are improved, and variations in material properties are decreased to 1/2. Even when the elongation of temper rolling is changed to 0.5 to 2%, and the elongation of the leveler is changed to 0 to 1%, strain age hardenability does not deteriorate.
  • Each of the resultant hot-rolled sheets was cold-rolled in the cold rolling step comprising pickling and cold rolling under the conditions shown in Table 16 to form a cold-rolled sheet.
  • Each of the thus-obtained cold-rolled sheet was box-annealed and then continuously annealed under the conditions shown in Table 16. Some of the cold-rolled sheets were temper-rolled after the cold-rolled sheet annealing step. Box annealing may not be carried out. In all cases, the annealing temperature of box annealing was the recrystallization temperature or more.
  • the thus-obtained cold-rolled annealed sheets were examined with respect to the amount of dissolved N, the microstructure, tensile properties, the r value, strain age hardenability, and the aging property.
  • the surfaces of the steel sheets of Nos. 17 and 18 were coated by hot-dip galvanization in an in-line after continuous annealing shown in Table 16 to form coated steel sheets.
  • the coated steel sheets were also examined with respect to the same properties as described above.
  • All the examples of the present invention have excellent ductility, extremely high stable BH amount and ⁇ TS, excellent strain age hardenability, a mean r value of as high as 1.2, and natural non-aging properties.
  • the properties of the hot-dip galvanized steel sheets of Nos. 17 and 18 shown in Table 17 are substantially the same as the cold-rolled steel sheets subjected to continuous annealing.
  • ductility deteriorates, the BH amount and TS are low, or aging deterioration significantly occurs. Therefore, the comparative examples do not have all the intended properties, and thus cannot be said as steel sheets having sufficient properties.
  • Steel sheet No. 11 contains C and N in amounts out of the range of the present invention, and has an amount of dissolved N and a martensite amount lower than the range of the present invention. Therefore, the BH amount and ⁇ TS are decreased, and ⁇ Y-E1 is increased.
  • Steel sheet No. 12 contains Al, N/Al and N out of the range of the present invention, and has an amount of dissolved N lower than the range of the present invention, and the average crystal grain diameter of ferrite higher than the range of the present invention. Therefore, the BH amount and ⁇ TS are decreased, and ⁇ Y-E1 is increased.
  • the slab heating temperature and finisher delivery temperature FDT are out of the range of the present invention, the amount of dissolved N and the amount of martensite are lower than the range of the present invention, and the average crystal grain diameter of ferrite is higher than the range of the present invention. Therefore, the r value, the BH amount and ⁇ TS are decreased.
  • the coiling temperature after hot rolling is out of the range of the present invention, the amount of dissolved N and the amount of martensite are lower than the range of the present invention, and the average crystal grain diameter of ferrite is higher than the range of the present invention. Therefore, the r value, the BH amount and ⁇ TS are decreased.
  • the continuous annealing temperature is out of the range of the present invention, martensite is not formed, and the average crystal grain diameter of ferrite is higher than the range of the present invention. Therefore, the BH amount and ⁇ TS are decreased, and ⁇ Y-E1 is increased.
  • box annealing is not performed to fail to develop the desirable aggregation structure, deteriorating the r value. Also, the average crystal grain diameter of ferrite, and the area ratio of martensite are out of the range of the present invention.
  • Example 18 Steel having the composition shown in Table 18 was formed in a slab by the same method as Example 1, and then heated and' roughly rolled under the conditions shown in Table 19 to form a sheet bar having a thickness of 30 mm.
  • the sheet bar was hot-rolled by the hot rolling step comprising finish rolling under the conditions shown in Table 19 to form a hot-rolled sheet.
  • the adjacent sheet bars on the finisher entrance side after rough rolling were bonded together' by the melt welding method, and then continuously rolled.
  • the temperatures of the ends of the sheet bar were controlled in the width direction and the length direction by using an induction heating-type sheet bar edge heater and a sheet bar heater.
  • the thus-obtained hot-rolled sheet was cold-rolled by the cold rolling step comprising pickling and cold rolling under the conditions shown in Table 19 to form a cold-rolled sheet having a thickness of 1.6 mm. Then, the cold-rolled sheet was box-annealed and then continuously annealed by a continuous annealing furnace under the conditions shown in Table 19. In all cases, the annealing temperature of box annealing are the recrystallization temperature or more.
  • the thus-obtained cold-rolled annealed sheet was examined with respect to the amount of dissolved N, the microstructure, tensile properties, the r value, and strain age hardenability by the same methods as Example 1.
  • All the examples of the present invention have excellent strain age hardenability and a high r value, and exhibit extremely high stable BH amount, ⁇ TS and mean r value regardless of variations in production conditions. It was also recognized that in the examples of the present invention, by performing continuous rolling and controlling the temperature of the sheet bar in the long direction and the width direction, the thickness precision and the shape of the produced steel sheet are improved, and variations in material properties are decreased.
  • a cold rolled steel sheet can be obtained, in which TS is greatly increased by press forming and heat treatment while maintaining excellent deep drawability in press forming.
  • the cold-rolled steel sheet has the excellent effect of industrially producing coated steel sheets by electro-galvanization, hot-dip galvanization, alloying hot-dip galvanization.
  • Heat treatment temperature (°C) 120 200 300 Retention time (min.) 2 82 114 133 5 86 119 136 10 91 122 138 20 95 125 140 Al % N/Al TS ⁇ r value MPa ⁇ TS MPa 0.020 0.75 775 58 0.036 0.42 762 55 0.049 0.31 753 42 0.072 0.21 720 25 0.080 0.19 719 19

Landscapes

  • Chemical & Material Sciences (AREA)
  • Engineering & Computer Science (AREA)
  • Mechanical Engineering (AREA)
  • Materials Engineering (AREA)
  • Metallurgy (AREA)
  • Organic Chemistry (AREA)
  • Physics & Mathematics (AREA)
  • Thermal Sciences (AREA)
  • Crystallography & Structural Chemistry (AREA)
  • Heat Treatment Of Sheet Steel (AREA)
  • Coating With Molten Metal (AREA)
EP01906128A 2000-05-26 2001-02-14 Kaltgewalztes stahlblech und galvanisiertes stahlblech mit guten reckalterungseigenschaften und herstellungsverfahren dafür Expired - Lifetime EP1291448B1 (de)

Priority Applications (2)

Application Number Priority Date Filing Date Title
EP04023101A EP1498507B1 (de) 2000-05-26 2001-02-14 Kaltgewalztes Stahlblech und Zinkblech mit Reckalterungseigenschaften und Verfahren zur dessen Herstellung
EP04023082A EP1498506B1 (de) 2000-05-26 2001-02-14 Hochfestes Kaltgewalztes Stahlblech mit hoch r-Wert, exzellenter Reckalterungseigenschaften und Alterungsbeständigkeit sowie Verfahren zur dessen Herstellung

Applications Claiming Priority (9)

Application Number Priority Date Filing Date Title
JP2000156274A JP4524859B2 (ja) 2000-05-26 2000-05-26 歪時効硬化特性に優れた深絞り用冷延鋼板およびその製造方法
JP2000156274 2000-05-26
JP2000193717 2000-06-28
JP2000193717 2000-06-28
JP2000328924 2000-10-27
JP2000328924 2000-10-27
JP2000335803A JP4665302B2 (ja) 2000-11-02 2000-11-02 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およびその製造方法
JP2000335803 2000-11-02
PCT/JP2001/001004 WO2001090431A1 (fr) 2000-05-26 2001-02-14 Tole d'acier laminee a froid et tole d'acier galvanisee possedant des proprietes de durcissement par ecrouissage et par precipitation et procede de production associe

Related Child Applications (2)

Application Number Title Priority Date Filing Date
EP04023101A Division EP1498507B1 (de) 2000-05-26 2001-02-14 Kaltgewalztes Stahlblech und Zinkblech mit Reckalterungseigenschaften und Verfahren zur dessen Herstellung
EP04023082A Division EP1498506B1 (de) 2000-05-26 2001-02-14 Hochfestes Kaltgewalztes Stahlblech mit hoch r-Wert, exzellenter Reckalterungseigenschaften und Alterungsbeständigkeit sowie Verfahren zur dessen Herstellung

Publications (3)

Publication Number Publication Date
EP1291448A1 true EP1291448A1 (de) 2003-03-12
EP1291448A4 EP1291448A4 (de) 2004-06-30
EP1291448B1 EP1291448B1 (de) 2006-06-28

Family

ID=27481319

Family Applications (3)

Application Number Title Priority Date Filing Date
EP01906128A Expired - Lifetime EP1291448B1 (de) 2000-05-26 2001-02-14 Kaltgewalztes stahlblech und galvanisiertes stahlblech mit guten reckalterungseigenschaften und herstellungsverfahren dafür
EP04023082A Expired - Lifetime EP1498506B1 (de) 2000-05-26 2001-02-14 Hochfestes Kaltgewalztes Stahlblech mit hoch r-Wert, exzellenter Reckalterungseigenschaften und Alterungsbeständigkeit sowie Verfahren zur dessen Herstellung
EP04023101A Expired - Lifetime EP1498507B1 (de) 2000-05-26 2001-02-14 Kaltgewalztes Stahlblech und Zinkblech mit Reckalterungseigenschaften und Verfahren zur dessen Herstellung

Family Applications After (2)

Application Number Title Priority Date Filing Date
EP04023082A Expired - Lifetime EP1498506B1 (de) 2000-05-26 2001-02-14 Hochfestes Kaltgewalztes Stahlblech mit hoch r-Wert, exzellenter Reckalterungseigenschaften und Alterungsbeständigkeit sowie Verfahren zur dessen Herstellung
EP04023101A Expired - Lifetime EP1498507B1 (de) 2000-05-26 2001-02-14 Kaltgewalztes Stahlblech und Zinkblech mit Reckalterungseigenschaften und Verfahren zur dessen Herstellung

Country Status (7)

Country Link
EP (3) EP1291448B1 (de)
KR (1) KR20020019124A (de)
CN (1) CN1158398C (de)
CA (1) CA2379698C (de)
DE (3) DE60121234T2 (de)
TW (1) TW565621B (de)
WO (1) WO2001090431A1 (de)

Cited By (27)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
EP1607490A1 (de) * 2004-06-18 2005-12-21 Nippon Steel Corporation Stahlblech geeignet als Weissblech mit ausgezeichneter Umformbarkeit und Verfahren zur dessen Herstellung.
WO2006026982A1 (de) * 2004-09-09 2006-03-16 Salzgitter Flachstahl Gmbh Beruhigter, unlegierter oder mikrolegierter walzstahl mit bake-hardening effekt und verfahren zu seiner herstellung
US7442268B2 (en) 2004-11-24 2008-10-28 Nucor Corporation Method of manufacturing cold rolled dual-phase steel sheet
US7608155B2 (en) 2006-09-27 2009-10-27 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
US7959747B2 (en) 2004-11-24 2011-06-14 Nucor Corporation Method of making cold rolled dual phase steel sheet
US8002016B2 (en) 2008-03-19 2011-08-23 Nucor Corporation Strip casting apparatus with casting roll positioning
US8337643B2 (en) 2004-11-24 2012-12-25 Nucor Corporation Hot rolled dual phase steel sheet
US8435363B2 (en) 2007-10-10 2013-05-07 Nucor Corporation Complex metallographic structured high strength steel and manufacturing same
US8631853B2 (en) 2008-03-19 2014-01-21 Nucor Corporation Strip casting apparatus for rapid set and change of casting rolls
RU2514743C2 (ru) * 2011-09-30 2014-05-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочной стальной лист, обладающий превосходной способностью к термическому упрочнению и формуемостью, и способ его производства
RU2518852C1 (ru) * 2012-07-20 2014-06-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист и способ его изготовления
RU2524031C2 (ru) * 2011-09-27 2014-07-27 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочные холоднокатаные стальные листы, обладающие превосходным качеством поверхности после штамповки, и способы их производства
US20140220383A1 (en) * 2011-08-05 2014-08-07 Jfe Steel Corporation Hot-dip galvanized steel sheet and production method therefor
RU2525013C1 (ru) * 2012-04-27 2014-08-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист, пригодный для химической конверсионной обработки, и способ его изготовления
RU2526345C2 (ru) * 2012-03-30 2014-08-20 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист, обладающий превосходной сгибаемостью, и способ его производства
RU2530212C2 (ru) * 2011-09-15 2014-10-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист и стальной лист с покрытием, обладающие превосходной способностью к термическому упрочнению и формуемостью, и способ их производства
RU2532782C2 (ru) * 2011-09-27 2014-11-10 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист, обладающий превосходным качеством поверхности после штамповки и способностью к упрочнению при обжиге, а также способ его производства
WO2015009416A1 (en) * 2013-03-15 2015-01-22 Thyssenkrupp Steel Usa, Llc High strength-high ductility cold rolled recovery annealed steel
EP2589677A4 (de) * 2010-06-29 2015-07-01 Jfe Steel Corp Hochfestes feuerverzinktes stahlblech mit ausgezeichneter verarbeitbarkeit und verfahren zur herstellung davon
US9200352B2 (en) * 2008-08-05 2015-12-01 Jfe Steel Corporation High strength galvannealed steel sheet with excellent appearance and method for manufacturing the same
EP2762583A4 (de) * 2011-09-30 2015-12-02 Nippon Steel & Sumitomo Metal Corp Hochfestes feuerverzinktes stahlblech mit hervorragender beständigkeit gegen verzögerte fraktur und herstellungsverfahren dafür
RU2571667C2 (ru) * 2012-03-30 2015-12-20 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист с превосходной стойкостью к старению и способ его изготовления
EP2128289B1 (de) 2007-02-28 2016-08-10 JFE Steel Corporation Stahlblech für dosen, heissgewalztes stahlblech zur verwendung als basismetall und herstellungsverfahren für beide
US10053749B2 (en) 2008-12-26 2018-08-21 Posco Production method for plated steel sheet using a steel sheet annealing device
RU2690076C1 (ru) * 2018-12-18 2019-05-30 Публичное акционерное общество "Северсталь" Листовой прокат и способ его получения
US11155902B2 (en) 2006-09-27 2021-10-26 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
EP4089188A4 (de) * 2020-01-08 2022-12-28 Nippon Steel Corporation Stahlblech und verfahren zur herstellung davon

Families Citing this family (36)

* 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
JP4635525B2 (ja) * 2003-09-26 2011-02-23 Ｊｆｅスチール株式会社 深絞り性に優れた高強度鋼板およびその製造方法
CA2552963C (en) 2004-01-14 2010-11-16 Nippon Steel Corporation Hot dip galvanized high strength steel sheet excellent in plating adhesion and hole expandability and method of production of same
CN102242306B (zh) 2005-08-03 2013-03-27 住友金属工业株式会社 热轧钢板及冷轧钢板及它们的制造方法
KR100868457B1 (ko) * 2007-05-31 2008-11-11 주식회사 포스코 도금밀착성이 우수한 합금화용융아연도금강판과 그제조방법
JP4712882B2 (ja) * 2008-07-11 2011-06-29 株式会社神戸製鋼所 耐水素脆化特性および加工性に優れた高強度冷延鋼板
JP2010126757A (ja) * 2008-11-27 2010-06-10 Jfe Steel Corp 高強度溶融亜鉛めっき鋼板およびその製造方法
JP5051247B2 (ja) * 2010-01-15 2012-10-17 Ｊｆｅスチール株式会社 成形性と形状凍結性に優れた冷延鋼板およびその製造方法
TWI475114B (zh) * 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 (ja) * 2010-11-05 2015-12-02 Ｊｆｅスチール株式会社 深絞り性および焼付硬化性に優れる高強度冷延鋼板とその製造方法
CN102041444A (zh) * 2010-12-21 2011-05-04 南阳汉冶特钢有限公司 一种低碳低硅优质碳素结构钢及其生产方法
TWI468530B (zh) 2012-02-13 2015-01-11 新日鐵住金股份有限公司 冷軋鋼板、鍍敷鋼板、及其等之製造方法
WO2013132821A1 (ja) * 2012-03-06 2013-09-12 Ｊｆｅスチール株式会社 温間プレス成形方法および自動車骨格部品
ES2718904T3 (es) 2012-04-13 2019-07-05 Nippon Steel Corp Lámina de acero para galvanoplastia, lámina de acero galvanizado, y métodos para la producción de la misma
KR101412293B1 (ko) * 2012-04-25 2014-06-25 현대제철 주식회사 고강도 냉연강판 및 그 제조 방법
KR101294447B1 (ko) * 2013-03-26 2013-08-07 현대제철 주식회사 성형성이 우수한 법랑용 냉연강판 제조 방법
KR101294577B1 (ko) * 2013-03-26 2013-08-07 현대제철 주식회사 딥드로잉 특성이 우수한 법랑용 냉연강판 제조 방법
KR101294477B1 (ko) * 2013-03-26 2013-08-07 현대제철 주식회사 딥드로잉성 및 밀착성이 우수한 법랑용 강판 제조 방법
KR101294575B1 (ko) * 2013-03-26 2013-08-07 현대제철 주식회사 내피쉬스케일성이 우수한 법랑용 냉연강판 제조 방법
JP5633594B2 (ja) * 2013-04-02 2014-12-03 Ｊｆｅスチール株式会社 打ち抜き性および耐熱ひずみ特性に優れた冷延鋼板およびその製造方法
EP2985361B1 (de) * 2013-10-02 2018-03-14 Nippon Steel & Sumitomo Metal Corporation Aushärtbarer stahl
CN103590352B (zh) * 2013-11-06 2015-07-08 湖北秭归新亚交通设施有限公司 废旧热浸锌公路钢护栏处理工艺
CA2932854C (en) * 2013-12-19 2017-12-12 Nisshin Steel Co., Ltd. Steel sheet hot-dip coated with zn-al-mg-based system having excellent workability and method for manufacturing same
KR101657845B1 (ko) * 2014-12-26 2016-09-20 주식회사 포스코 박슬라브 표면 품질이 우수한 고강도 냉연강판 및 그 제조방법
CN104950777B (zh) * 2015-06-13 2017-11-07 合肥荣事达电子电器集团有限公司 一种日夜通用的室外监控装置
CN106244914A (zh) * 2015-06-13 2016-12-21 郭策 室外监控装置的制备方法
KR102031449B1 (ko) * 2017-12-24 2019-10-11 주식회사 포스코 상온내시효성 및 소부경화성이 우수한 아연계 도금강판 및 그 제조방법
RU2723872C1 (ru) * 2019-05-23 2020-06-17 Публичное акционерное общество "Магнитогорский металлургический комбинат" Способ производства холоднокатаного высокопрочного листового проката из низколегированной стали
KR102201438B1 (ko) * 2019-06-26 2021-01-12 현대제철 주식회사 강판 및 그 제조방법
CN111663075B (zh) * 2020-04-09 2021-10-01 北京首钢股份有限公司 一种冲压用酸洗钢及其制备方法、应用
CN111926252B (zh) * 2020-07-31 2022-01-18 马鞍山钢铁股份有限公司 一种深冲用途的热轧酸洗钢板及其生产方法
KR102426248B1 (ko) * 2020-11-05 2022-07-28 주식회사 포스코 선영성이 우수한 고강도 아연계 도금강판 및 그 제조방법
KR102468036B1 (ko) * 2020-11-12 2022-11-17 주식회사 포스코 성형성이 우수한 고강도 아연계 도금강판 및 그 제조방법
KR20220086058A (ko) * 2020-12-16 2022-06-23 주식회사 포스코 재질 균일성이 우수한 고강도 냉연, 도금 강판 및 이들의 제조 방법
TWI768666B (zh) * 2021-01-20 2022-06-21 中國鋼鐵股份有限公司 高成形性冷軋鋼材及其製造方法
CN113462963B (zh) * 2021-06-10 2022-05-20 江阴市万众精密机械有限公司 一种增速箱联轴器用耐冲击、耐低温止推盘及其制备方法

Citations (19)

* 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
JPS6052528A (ja) * 1983-09-02 1985-03-25 Kawasaki Steel Corp 延性およびスポツト溶接性の良好な高強度薄鋼板の製造方法
JPS60145355A (ja) * 1984-01-06 1985-07-31 Kawasaki Steel Corp 延性が良好で時効劣化のない低降伏比高張力熱延鋼板とその製造方法
JPS61272323A (ja) * 1985-05-28 1986-12-02 Kawasaki Steel Corp 連続焼鈍による表面処理用原板の製造方法
US4838955A (en) * 1985-07-29 1989-06-13 Hoogovens Groep B.V. Method for the manufacture of hard steel sheet from Al-killed continuous-cast carbon-manganese steel
EP0429094A1 (de) * 1984-01-20 1991-05-29 KABUSHIKI KAISHA KOBE SEIKO SHO also known as Kobe Steel Ltd. Hochfester, niedriggekohlter Stahl, Gegenstände daraus und Verfahren zur Herstellung dieses Stahls
JPH0474824A (ja) * 1990-07-18 1992-03-10 Sumitomo Metal Ind Ltd 焼付硬化性と加工性に優れた熱延鋼板の製造方法
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
EP0510718A2 (de) * 1991-04-26 1992-10-28 Kawasaki Steel Corporation Hochfestes, kaltgewalztes, bei Raumtemperatur alterungsbeständiges, tiefziehbares Stahlblech und Herstellungsverfahren
EP0608430A1 (de) * 1992-06-22 1994-08-03 Nippon Steel Corporation Kaltgewalzte stahlplatte mit guter einbrennhärtbarkeit, und formbarbeit ohne kaltalterungserscheinungen und gegossene zinkbeschichtete kaltgewalzte stahlplatte und herstellungsverfahren
EP0612857A1 (de) * 1992-09-14 1994-08-31 Nippon Steel Corporation Ferritisch einphasige stahlplatte oder zinkplattierte stahlplatte zum tiefziehen ohne kaltalternungsenscheinungen und verfahren zu dessen herstellung
JPH0790482A (ja) * 1993-09-21 1995-04-04 Kawasaki Steel Corp 耐衝撃性に優れた薄鋼板およびその製造方法
US5405463A (en) * 1980-10-24 1995-04-11 Nippon Kokan Kabushiki Kaisha Continuous annealing process of producing cold rolled mild steel sheet excellent in deep drawability and aging resistibility
EP0659890A2 (de) * 1993-12-21 1995-06-28 Kawasaki Steel Corporation Verfahren zum Herstellen von dünnen Stahlblechern mit niedriger planarer Anisotropie für Dosen
US5582658A (en) * 1990-08-17 1996-12-10 Kawasaki Steel Corporation High strength steel sheet adapted for press forming and method of producing the same
EP0769565A1 (de) * 1995-03-27 1997-04-23 Nippon Steel Corporation Kaltgewalztes blech mit extrem niedrigem kohlenstoffgehalt und galvanisiertes blech, beide mit hervorragenden ermündungseigenschaften und verfahren zu deren herstellung
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
EP0943696A1 (de) * 1997-09-04 1999-09-22 Kawasaki Steel Corporation Stahlplatten für fässer, verfahren zu deren herstellung und fässer
EP1002884A1 (de) * 1998-04-27 2000-05-24 Nkk Corporation Kaltgewalzte stahlplatte exzellenter formbarkeit, flachförmigen eigenschaften und eindellwiderstand, feuerverzinkte stahlplatte und verfahren zur deren herstellung

Family Cites Families (6)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
JPS5842753A (ja) * 1981-09-07 1983-03-12 Kobe Steel Ltd 高ｒ値型複合組織高強度冷延鋼板及びその製造方法
JP3303931B2 (ja) * 1992-10-06 2002-07-22 川崎製鉄株式会社 焼付け硬化性を有する高強度缶用薄鋼板及びその製造方法
JP3383017B2 (ja) * 1993-07-19 2003-03-04 新日本製鐵株式会社 加工性に優れた焼付け硬化性高強度冷延鋼板の製造方法
JP3390256B2 (ja) * 1994-07-21 2003-03-24 川崎製鉄株式会社 焼付け硬化性及び耐時効性に優れた高強度高加工性製缶用鋼板及びその製造方法
JP3527092B2 (ja) * 1998-03-27 2004-05-17 新日本製鐵株式会社 加工性の良い高強度合金化溶融亜鉛めっき鋼板とその製造方法
JP6102816B2 (ja) * 2014-03-31 2017-03-29 ブラザー工業株式会社 成膜装置、成膜方法及び成膜プログラム

Patent Citations (19)

* 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
US5405463A (en) * 1980-10-24 1995-04-11 Nippon Kokan Kabushiki Kaisha Continuous annealing process of producing cold rolled mild steel sheet excellent in deep drawability and aging resistibility
JPS6052528A (ja) * 1983-09-02 1985-03-25 Kawasaki Steel Corp 延性およびスポツト溶接性の良好な高強度薄鋼板の製造方法
JPS60145355A (ja) * 1984-01-06 1985-07-31 Kawasaki Steel Corp 延性が良好で時効劣化のない低降伏比高張力熱延鋼板とその製造方法
EP0429094A1 (de) * 1984-01-20 1991-05-29 KABUSHIKI KAISHA KOBE SEIKO SHO also known as Kobe Steel Ltd. Hochfester, niedriggekohlter Stahl, Gegenstände daraus und Verfahren zur Herstellung dieses Stahls
JPS61272323A (ja) * 1985-05-28 1986-12-02 Kawasaki Steel Corp 連続焼鈍による表面処理用原板の製造方法
US4838955A (en) * 1985-07-29 1989-06-13 Hoogovens Groep B.V. Method for the manufacture of hard steel sheet from Al-killed continuous-cast carbon-manganese steel
JPH0474824A (ja) * 1990-07-18 1992-03-10 Sumitomo Metal Ind Ltd 焼付硬化性と加工性に優れた熱延鋼板の製造方法
US5582658A (en) * 1990-08-17 1996-12-10 Kawasaki 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
EP0510718A2 (de) * 1991-04-26 1992-10-28 Kawasaki Steel Corporation Hochfestes, kaltgewalztes, bei Raumtemperatur alterungsbeständiges, tiefziehbares Stahlblech und Herstellungsverfahren
EP0608430A1 (de) * 1992-06-22 1994-08-03 Nippon Steel Corporation Kaltgewalzte stahlplatte mit guter einbrennhärtbarkeit, und formbarbeit ohne kaltalterungserscheinungen und gegossene zinkbeschichtete kaltgewalzte stahlplatte und herstellungsverfahren
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
EP0612857A1 (de) * 1992-09-14 1994-08-31 Nippon Steel Corporation Ferritisch einphasige stahlplatte oder zinkplattierte stahlplatte zum tiefziehen ohne kaltalternungsenscheinungen und verfahren zu dessen herstellung
JPH0790482A (ja) * 1993-09-21 1995-04-04 Kawasaki Steel Corp 耐衝撃性に優れた薄鋼板およびその製造方法
EP0659890A2 (de) * 1993-12-21 1995-06-28 Kawasaki Steel Corporation Verfahren zum Herstellen von dünnen Stahlblechern mit niedriger planarer Anisotropie für Dosen
EP0769565A1 (de) * 1995-03-27 1997-04-23 Nippon Steel Corporation Kaltgewalztes blech mit extrem niedrigem kohlenstoffgehalt und galvanisiertes blech, beide mit hervorragenden ermündungseigenschaften und verfahren zu deren herstellung
EP0943696A1 (de) * 1997-09-04 1999-09-22 Kawasaki Steel Corporation Stahlplatten für fässer, verfahren zu deren herstellung und fässer
EP1002884A1 (de) * 1998-04-27 2000-05-24 Nkk Corporation Kaltgewalzte stahlplatte exzellenter formbarkeit, flachförmigen eigenschaften und eindellwiderstand, feuerverzinkte stahlplatte und verfahren zur deren herstellung

Non-Patent Citations (7)

* Cited by examiner, † Cited by third party
Title
PATENT ABSTRACTS OF JAPAN vol. 005, no. 012 (C-040), 24 January 1981 (1981-01-24) -& JP 55 141526 A (KAWASAKI STEEL CORP), 5 November 1980 (1980-11-05) *
PATENT ABSTRACTS OF JAPAN vol. 009, no. 184 (C-294), 30 July 1985 (1985-07-30) -& JP 60 052528 A (KAWASAKI SEITETSU KK), 25 March 1985 (1985-03-25) *
PATENT ABSTRACTS OF JAPAN vol. 009, no. 306 (C-317), 3 December 1985 (1985-12-03) -& JP 60 145355 A (KAWASAKI SEITETSU KK), 31 July 1985 (1985-07-31) *
PATENT ABSTRACTS OF JAPAN vol. 011, no. 136 (C-419), 30 April 1987 (1987-04-30) -& JP 61 272323 A (KAWASAKI STEEL CORP), 2 December 1986 (1986-12-02) *
PATENT ABSTRACTS OF JAPAN vol. 016, no. 283 (C-0955), 24 June 1992 (1992-06-24) -& JP 04 074824 A (SUMITOMO METAL IND LTD), 10 March 1992 (1992-03-10) *
PATENT ABSTRACTS OF JAPAN vol. 1995, no. 07, 31 August 1995 (1995-08-31) -& JP 07 090482 A (KAWASAKI STEEL CORP), 4 April 1995 (1995-04-04) *
See also references of WO0190431A1 *

Cited By (36)

* Cited by examiner, † Cited by third party
Publication number Priority date Publication date Assignee Title
US8012276B2 (en) 2004-06-18 2011-09-06 Nippon Steel Corporation Method for manufacturing a steel sheet for tin plated steel sheet and tin-free steel sheet each having excellent formability
EP1607490A1 (de) * 2004-06-18 2005-12-21 Nippon Steel Corporation Stahlblech geeignet als Weissblech mit ausgezeichneter Umformbarkeit und Verfahren zur dessen Herstellung.
US7501031B2 (en) 2004-06-18 2009-03-10 Nippon Steel Corporation Steel sheet for tin plated steel sheet and tin-free steel sheet each having excellent formability and manufacturing method thereof
WO2006026982A1 (de) * 2004-09-09 2006-03-16 Salzgitter Flachstahl Gmbh Beruhigter, unlegierter oder mikrolegierter walzstahl mit bake-hardening effekt und verfahren zu seiner herstellung
US7959747B2 (en) 2004-11-24 2011-06-14 Nucor Corporation Method of making cold rolled dual phase steel sheet
US7879160B2 (en) 2004-11-24 2011-02-01 Nucor Corporation Cold rolled dual-phase steel sheet
US8337643B2 (en) 2004-11-24 2012-12-25 Nucor Corporation Hot rolled dual phase steel sheet
US8366844B2 (en) 2004-11-24 2013-02-05 Nucor Corporation Method of making 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
US7608155B2 (en) 2006-09-27 2009-10-27 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
US11155902B2 (en) 2006-09-27 2021-10-26 Nucor Corporation High strength, hot dip coated, dual phase, steel sheet and method of manufacturing same
EP2128289B1 (de) 2007-02-28 2016-08-10 JFE Steel Corporation Stahlblech für dosen, heissgewalztes stahlblech zur verwendung als basismetall und herstellungsverfahren für beide
EP2128289B2 (de) † 2007-02-28 2019-10-23 JFE Steel Corporation Stahlblech für dosen, heissgewalztes stahlblech zur verwendung als basismetall und herstellungsverfahren für beide
US9157138B2 (en) 2007-10-10 2015-10-13 Nucor Corporation Complex metallographic structured high strength steel and method of manufacturing
US8435363B2 (en) 2007-10-10 2013-05-07 Nucor Corporation Complex metallographic structured high strength steel and manufacturing same
US8631853B2 (en) 2008-03-19 2014-01-21 Nucor Corporation Strip casting apparatus for rapid set and change of casting rolls
US8875777B2 (en) 2008-03-19 2014-11-04 Nucor Corporation Strip casting apparatus for rapid set and change of casting rolls
US9120147B2 (en) 2008-03-19 2015-09-01 Nucor Corporation Strip casting apparatus for rapid set and change of casting rolls
US8002016B2 (en) 2008-03-19 2011-08-23 Nucor Corporation Strip casting apparatus with casting roll positioning
US9200352B2 (en) * 2008-08-05 2015-12-01 Jfe Steel Corporation High strength galvannealed steel sheet with excellent appearance and method for manufacturing the same
US10053749B2 (en) 2008-12-26 2018-08-21 Posco Production method for plated steel sheet using a steel sheet annealing device
EP2589677A4 (de) * 2010-06-29 2015-07-01 Jfe Steel Corp Hochfestes feuerverzinktes stahlblech mit ausgezeichneter verarbeitbarkeit und verfahren zur herstellung davon
US20140220383A1 (en) * 2011-08-05 2014-08-07 Jfe Steel Corporation Hot-dip galvanized steel sheet and production method therefor
US10337094B2 (en) * 2011-08-05 2019-07-02 Jfe Steel Corporation Hot-dip galvanized steel sheet and production method therefor
RU2530212C2 (ru) * 2011-09-15 2014-10-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист и стальной лист с покрытием, обладающие превосходной способностью к термическому упрочнению и формуемостью, и способ их производства
RU2532782C2 (ru) * 2011-09-27 2014-11-10 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист, обладающий превосходным качеством поверхности после штамповки и способностью к упрочнению при обжиге, а также способ его производства
RU2524031C2 (ru) * 2011-09-27 2014-07-27 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочные холоднокатаные стальные листы, обладающие превосходным качеством поверхности после штамповки, и способы их производства
EP2762583A4 (de) * 2011-09-30 2015-12-02 Nippon Steel & Sumitomo Metal Corp Hochfestes feuerverzinktes stahlblech mit hervorragender beständigkeit gegen verzögerte fraktur und herstellungsverfahren dafür
RU2514743C2 (ru) * 2011-09-30 2014-05-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочной стальной лист, обладающий превосходной способностью к термическому упрочнению и формуемостью, и способ его производства
RU2571667C2 (ru) * 2012-03-30 2015-12-20 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист с превосходной стойкостью к старению и способ его изготовления
RU2526345C2 (ru) * 2012-03-30 2014-08-20 ДжФЕ СТИЛ КОРПОРЕЙШН Холоднокатаный стальной лист, обладающий превосходной сгибаемостью, и способ его производства
RU2525013C1 (ru) * 2012-04-27 2014-08-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист, пригодный для химической конверсионной обработки, и способ его изготовления
RU2518852C1 (ru) * 2012-07-20 2014-06-10 ДжФЕ СТИЛ КОРПОРЕЙШН Высокопрочный холоднокатаный стальной лист и способ его изготовления
WO2015009416A1 (en) * 2013-03-15 2015-01-22 Thyssenkrupp Steel Usa, Llc High strength-high ductility cold rolled recovery annealed steel
RU2690076C1 (ru) * 2018-12-18 2019-05-30 Публичное акционерное общество "Северсталь" Листовой прокат и способ его получения
EP4089188A4 (de) * 2020-01-08 2022-12-28 Nippon Steel Corporation Stahlblech und verfahren zur herstellung davon

Also Published As

Publication number Publication date
TW565621B (en) 2003-12-11
CA2379698C (en) 2009-02-17
CN1386140A (zh) 2002-12-18
EP1498507A1 (de) 2005-01-19
DE60121234D1 (de) 2006-08-10
EP1498507B1 (de) 2006-06-28
DE60121234T2 (de) 2006-11-09
DE60121162T2 (de) 2006-11-09
DE60121233D1 (de) 2006-08-10
EP1498506B1 (de) 2006-06-28
EP1291448A4 (de) 2004-06-30
DE60121233T2 (de) 2006-11-09
EP1498506A1 (de) 2005-01-19
EP1291448B1 (de) 2006-06-28
CN1158398C (zh) 2004-07-21
KR20020019124A (ko) 2002-03-09
WO2001090431A1 (fr) 2001-11-29
CA2379698A1 (en) 2001-11-29
DE60121162D1 (de) 2006-08-10

Similar Documents

Publication Publication Date Title
US7067023B2 (en) Cold rolled steel sheet and galvanized steel sheet having strain age hardenability and method of producing the same
EP1498506B1 (de) Hochfestes Kaltgewalztes Stahlblech mit hoch r-Wert, exzellenter Reckalterungseigenschaften und Alterungsbeständigkeit sowie Verfahren zur dessen Herstellung
EP3309273B1 (de) Galvanisiertes stahlblech und verfahren zu dessen herstellung
US10662495B2 (en) High-strength steel sheet and production method for same, and production method for high-strength galvanized steel sheet
US10662496B2 (en) High-strength steel sheet and production method for same, and production method for high-strength galvanized steel sheet
KR101264574B1 (ko) 딥 드로잉성이 우수한 고강도 강판의 제조 방법
EP1191114A1 (de) Hochfestes warmgewalztes stahlblech mit ausgezeichneten reckalterungseigenschaften und herstellungsverfahren dafür
EP3187601A1 (de) Hochfestes stahlblech und verfahren zur herstellung davon
JP3846206B2 (ja) 歪時効硬化特性に優れた高張力冷延鋼板およびその製造方法
EP3178953A1 (de) Hochfestes stahlblech und herstellungsverfahren dafür und herstellungsverfahren für hochfestes verzinktes stahlblech
EP2623622A1 (de) Hochfestes feuerverzinktes stahlblech mit ausgezeichneter tiefziehbarkeit und dehnbarkeit sowie verfahren zu seiner herstellung
KR20210094620A (ko) 고강도 용융 아연 도금 강판 및 그의 제조 방법
JP6769576B1 (ja) 高強度溶融亜鉛めっき鋼板およびその製造方法
JP4752522B2 (ja) 深絞り用高強度複合組織型冷延鋼板の製造方法
CN115461482B (zh) 钢板、部件及其制造方法
JP2002241897A (ja) 降伏強さと破断伸びの変動が小さく高成形性と低降伏比とを有する鋼板およびその製造方法
JP4665302B2 (ja) 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およびその製造方法
WO2020136990A1 (ja) 高強度溶融亜鉛めっき鋼板およびその製造方法
JP4519373B2 (ja) 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およびその製造方法
WO2022054221A1 (ja) 鋼板およびその製造方法

Legal Events

Date Code Title Description
PUAI Public reference made under article 153(3) epc to a published international application that has entered the european phase

Free format text: ORIGINAL CODE: 0009012

17P Request for examination filed

Effective date: 20021219

AK Designated contracting states

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

Kind code of ref document: A1

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

RIC1 Information provided on ipc code assigned before grant

Ipc: 7C 22C 38/06 B

Ipc: 7C 22C 38/58 B

Ipc: 7C 22C 38/04 B

Ipc: 7C 21D 9/46 B

Ipc: 7C 21D 8/02 B

Ipc: 7C 22C 38/00 A

RAP1 Party data changed (applicant data changed or rights of an application transferred)

Owner name: JFE STEEL CORPORATION

RBV Designated contracting states (corrected)

Designated state(s): BE DE FR GB IT

A4 Supplementary search report drawn up and despatched

Effective date: 20040517

17Q First examination report despatched

Effective date: 20040714

GRAP Despatch of communication of intention to grant a patent

Free format text: ORIGINAL CODE: EPIDOSNIGR1

GRAS Grant fee paid

Free format text: ORIGINAL CODE: EPIDOSNIGR3

GRAA (expected) grant

Free format text: ORIGINAL CODE: 0009210

AK Designated contracting states

Kind code of ref document: B1

Designated state(s): BE DE FR GB IT

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: IT

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT;WARNING: LAPSES OF ITALIAN PATENTS WITH EFFECTIVE DATE BEFORE 2007 MAY HAVE OCCURRED AT ANY TIME BEFORE 2007. THE CORRECT EFFECTIVE DATE MAY BE DIFFERENT FROM THE ONE RECORDED.

Effective date: 20060628

Ref country code: BE

Free format text: LAPSE BECAUSE OF FAILURE TO SUBMIT A TRANSLATION OF THE DESCRIPTION OR TO PAY THE FEE WITHIN THE PRESCRIBED TIME-LIMIT

Effective date: 20060628

REG Reference to a national code

Ref country code: GB

Ref legal event code: FG4D

REF Corresponds to:

Ref document number: 60121162

Country of ref document: DE

Date of ref document: 20060810

Kind code of ref document: P

ET Fr: translation filed
PLBE No opposition filed within time limit

Free format text: ORIGINAL CODE: 0009261

STAA Information on the status of an ep patent application or granted ep patent

Free format text: STATUS: NO OPPOSITION FILED WITHIN TIME LIMIT

26N No opposition filed

Effective date: 20070329

GBPC Gb: european patent ceased through non-payment of renewal fee

Effective date: 20070214

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: GB

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20070214

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 16

REG Reference to a national code

Ref country code: FR

Ref legal event code: PLFP

Year of fee payment: 17

PGFP Annual fee paid to national office [announced via postgrant information from national office to epo]

Ref country code: FR

Payment date: 20170112

Year of fee payment: 17

Ref country code: DE

Payment date: 20170207

Year of fee payment: 17

REG Reference to a national code

Ref country code: DE

Ref legal event code: R119

Ref document number: 60121162

Country of ref document: DE

REG Reference to a national code

Ref country code: FR

Ref legal event code: ST

Effective date: 20181031

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: DE

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20180901

PG25 Lapsed in a contracting state [announced via postgrant information from national office to epo]

Ref country code: FR

Free format text: LAPSE BECAUSE OF NON-PAYMENT OF DUE FEES

Effective date: 20180228