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
  • C21D8/005—Modifying the physical properties by deformation combined with, or followed by, heat treatment of ferrous 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
  • C21D8/0226—Hot rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
  • C21D8/0236—Cold rolling
• • C—CHEMISTRY; METALLURGY
  • C21—METALLURGY OF IRON
  • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
  • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
  • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
  • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
  • C21D8/0273—Final recrystallisation annealing
• • 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
  • C22C38/008—Ferrous alloys, e.g. steel alloys containing tin
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/08—Ferrous alloys, e.g. steel alloys containing nickel
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/12—Ferrous alloys, e.g. steel alloys containing tungsten, tantalum, molybdenum, vanadium, or niobium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/20—Ferrous alloys, e.g. steel alloys containing chromium with copper
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/24—Ferrous alloys, e.g. steel alloys containing chromium with vanadium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/26—Ferrous alloys, e.g. steel alloys containing chromium with niobium or tantalum
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/28—Ferrous alloys, e.g. steel alloys containing chromium with titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/42—Ferrous alloys, e.g. steel alloys containing chromium with nickel with copper
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/46—Ferrous alloys, e.g. steel alloys containing chromium with nickel with vanadium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/48—Ferrous alloys, e.g. steel alloys containing chromium with nickel with niobium or tantalum
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
  • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
  • C22C38/50—Ferrous alloys, e.g. steel alloys containing chromium with nickel with titanium or zirconium
• • C—CHEMISTRY; METALLURGY
  • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
  • C22C—ALLOYS
  • C22C38/00—Ferrous alloys, e.g. steel alloys
  • C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
• • C—CHEMISTRY; METALLURGY
  • 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

Definitions

• the present invention relates to a high-strength cold rolled steel sheet suitable for use in an outer panel and the like of an automobile body and having a tensile strength TS of not less than 440 MPa, an average r-value of not less than 1.20 and a Bake hardening (BH) value of not less than 40 MPa and excellent deep drawability and bake hardenability as well as a method for manufacturing the same.
• TS tensile strength
• BH Bake hardening
• the effect of reducing the weight of the automobile body can be received as the strength of the steel sheet becomes higher.
• high-strength steel sheets having a tensile strength of not less than 440 MPa tend to be used in the automobile body recently.
• many members constituting the automobile body are formed by press working, so that the steel sheet as a raw material is required to have an excellent formability. Therefore, in order to attain the weight reduction and increase in strength of the automobile body, it is required to develop high-strength steel sheets having a tensile strength of not less than 440 MPa and an excellent deep drawability, concretely a Lankford value (r-value) indicating the deep drawability of not less than 1.2 as an average r-value.
• r-value Lankford value
• the strength after paint baking is high, and therefore it is also required that the bake hardenability (BH property) is excellent.
• the conventional steel sheet having an improved BH property has a tendency that the formability and deep drawability is poor as compared with the usual mild steel sheet because a greater amount of solute C is contained.
• the steel sheet used in the automobile body is required to have further excellent bake hardenability in addition to the high strength and excellent deep drawability.
• IF (interstitial free) steel is obtained by adding Ti and Nb to an extra low carbon steel to fix solute C and solute N and then a solid solution strengthening elements such as Si, Mn, P and so on are added thereto.
• Patent Document 1 discloses a high-tension cold rolled steel sheet having a chemical composition of C: 0.002-0.015%, Nb: (C ⁇ 3)% to (C ⁇ 8+0.020)%, Si: not more than 1.2%, Mn: 0.04-0.8% and P: 0.03-0.10% and possessing a non-aging property, a tensile strength of 35-45 kgf/mm 2 grade (340-440 MPa grade) and an excellent formability.
• it is required to add a greater amount of alloying elements for providing the tensile strength of not less than 440 MPa, so that there are problems that the r-value is lowered and the surface property and zinc coatability are deteriorated.
• the solute C and solute N are fixed by Ti and Nb, there are problems that secondary work brittleness is actualized and BH property effective for ensuring the dent resistance is not obtained.
• a dual phase steel sheet comprising a soft ferrite phase and a hard martensite phase generally has such characteristics that ductility is good and strength-ductility balance is excellent and yield ratio is low.
• the dual phase steel sheet has an excellent formability, whereas there is a problem that the deep drawability is poor because the r-value is low. This is considered due to the fact that martensite phase not contributing to the r-value in view of crystal orientation is present and the solute C required for the formation of martensite phase obstructs the formation of ⁇ 111 ⁇ recrystallization texture effective for increasing the r-value.
• Patent Document 2 proposes a technique that a dual phase steel sheet having a r-value of not less than 1.3 and a strength of 40-60 kgf/mm 2 is obtained by subjecting a steel raw material containing C: 0.05-0.15%, Si: not more than 1.50%, Mn: 0.30-1.50%, P: not more than 0.030%, S: not more than 0.030%, sol.
• Al 0.020-0.070% and N: 0.0020-0.0080% to hot rolling and cold rolling under predetermined conditions, conducting box annealing at a temperature ranging from recrystallization temperature to Ac 3 transformation point to precipitate AlN and enhance ⁇ 111 ⁇ texture, and then conducting temper rolling and further subjecting to continuous annealing at 700-800° C., a quenching and a tempering at 200-500° C.
• Patent Document 3 proposes a technique that a steel sheet having a ferrite-martensite dual phase and excellent deep drawability and shape fixability is obtained by subjecting a steel raw material containing C: not more than 0.20%, Si: not more than 1.0%, Mn: 0.8-2.5%, sol. Al: 0.01-0.20%, N: 0.0015-0.0150% and P: not more than 0.10% to hot rolling and cold rolling, subjecting the resulting steel sheet to box annealing at a temperature zone of 650-800° C. to thereby form recrystallization texture suitable for r-value and further segregate C and Mn atoms into austenite phase, and then subjecting to continuous annealing comprised of a step of heating above 600° C. and a cooling step.
• Patent Document 4 proposes a technique that a steel sheet. with a microstructure containing 3-100% in total of one or more of bainite, martensite and austenite and having an excellent deep drawability is obtained by subjecting a steel raw material containing, by mass, C: 0.03-0.25%, Si: 0.001-3.0%, Mn: 0.01-3.0%, P: 0.001-0.06%, S: not more than 0.05%, N: 0.001-0.030% and Al: 0.005-0.3% to hot rolling and cold rolling at a rolling reduction of not less than 30% but less than 95%, subjecting the resulting steel sheet to an annealing by heating at an average heating rate of 4-200° C./hr up to a maximum achieving temperature of 600-800° C. to thereby form, cluster or precipitate of Al and N for a desired texture, and further heating to a temperature of from Ac 1 transformation point to 1050° C. for ferrite-austenite dual phase zone and then cooling.
• Patent Documents 2-4 it is required to take an annealing step for enhancing the r-value by developing the texture through the formation of cluster or precipitation of Al and N and a heat-treating step for forming the desired microstructure. Further, the above annealing step is based on the box annealing and takes a long time because the holding time for soaking is not less than 1 hour. That is, the techniques of Patent Documents 2-4 take many step number in addition to the long annealing time, so that they are poor in the productivity.
• the annealing is conducted at a higher temperature over a long time of period at a coiled state, so that there are problems in quality that steel sheets are closely adhered to each other or temper color is caused, and a problem in production equipment that the service life of furnace body or inner cover in the annealing furnace is lowered.
• Patent Document 5 proposes a method for producing a dual phase steel sheet dispersed a given amount of a second phase (martensite and/or bainite) into ferrite by subjecting a steel raw material containing, by weight, C: 0.003-0.03%, Si: 0.2-1%, Mn: 0.3-1.5%, Al: 0.01-0.07% and Ti: 0.02-0.2% and having an atomic concentration ratio of (effective Ti)/(C+N) of 0.4-0.8 to hot rolling and cold rolling, and then subjecting the resulting steel sheet to continuous annealing comprised of a step of heating at a temperature of from Ac1 transformation point to 900° C. for 30 seconds to 10 minutes and a step of cooling at an average cooling rate of not less than 30° C./s.
• a second phase martensite and/or bainite
• a dual phase steel sheet having a r-value of 1.61 and a tensile strength of 482 MPa is obtained by subjecting a steel raw material having a chemical composition of C: 0.012%, Si: 0.32%, Mn: 0.53%, P: 0.03%, Al: 0.03% and Ti: 0.051% by weight to hot rolling and cold rolling and then to continuous annealing by annealing at 870° C. of ferrite-austenite dual phase zone for 2 minutes and quenching at an average cooling rate of 100° C./s.
• Patent Document 6 proposes a method of producing a high-tension cold rolled steel sheet of a dual phase type with a microstructure comprising ferrite phase as a main phase and not less than 1% as an area ratio of martensite phase and having an excellent deep drawability by subjecting a steel raw material containing, by mass, C: 0.01-0.08%, Si: not more than 2.0%, Mn: not more than 3.0%, Al: 0.005-0.20%, N: not more than 0.02% and V: 0.01-0.5% and satisfying a predetermined relation of V and C to hot rolling and cold rolling and subsequently to continuous annealing (recrystallization annealing) at a temperature zone of Ac 1 -Ac 3 transformation point.
• This technique is characterized in that the r-value is increased by rationalizing the V and C contents to thereby precipitate C in steel as a V carbide and reduce solute C as far as possible, and at the subsequent recrystallization annealing the steel sheet. is heated to the ferrite-austenite dual phase zone, whereby the V carbide is dissolved to incrassate C in austenite and then martensite is formed at the subsequent cooling step to increase the strength.
• Patent Document 7 proposes a method of producing a high-strength steel sheet by subjecting a steel raw material containing, by mass, C: 0.010-0.050%, Si: not more than 1.0%, Mn: 1.0-3.0%, P: 0.005-0.1%, S: not more than 0.01%, Al: 0.005-0.5%, N: not more than 0.01% and Nb: 0.01-0.3% and having Nb and C contents satisfying (Nb/93)/(C/12): 0.2-0.7 to hot rolling and cold rolling and then subjecting to annealing comprised of a step of heating to a ferrite-austenite dual phase temperature zone of 800-950° C. and a step of cooling at an average cooling rate of not less than 5° C./s within a temperature range of from the above annealing temperature to 500° C.
• Patent Document 7 is characterized in that the microstructure of the hot rolled steel sheet. is finely divided by the addition of Nb and further the Nb and C contents are controlled to (Nb/93)/(C/12): 0.2-0.7 to precipitate a part of C in steel during the hot rolling as NbC and solute C before the annealing, whereby the generation of ⁇ 111 ⁇ recrystallized grains is promoted from grain boundaries in the annealing to thereby increase the r-value, while martensite is produced by solute C not fixed as NbC in the cooling after the annealing to thereby increase the strength.
• Patent Document 7 there can be produced a high-strength steel sheet with a microstructure comprising a ferrite phase with an area ratio of not less than 50% and a martensite phase with an area ratio of not less than 1% and having an average r-value of not less than 1.2.
• Nb is a very expensive element, which is disadvantageous in the cost of the raw material. Also, Nb considerably delays the recrystallization of austenite, so that there is a problem that the load becomes high in the hot rolling. Further, NbC precipitated in the hot rolled steel sheet increases the deformation resistance in the cold rolling, so that when the cold rolling is carried out at a high rolling reduction (65%) as disclosed in examples of Patent Document 7, the rolling load becomes higher and hence the risk of causing troubles becomes large. Further the productivity decreases and the available steel sheet width becomes restricted.
• the technique of this document has many problems in view of stably producing the steel sheets.
• Patent Document 1 JP-A-556-139654
• Patent Document 2 JP-B-555-10650
• Patent Document 3 JP-A-555-100934
• Patent Document 4 JP-A-2003-64444
• Patent Document 5 JP-B-H01-35900
• Patent Document 6 JP-A-2002-226941
• Patent Document 7 JP-A-2005-120467
• the conventional techniques utilizing solid solution enhancement for increasing the strength of mild steel sheets having an excellent deep drawability are necessary to add a greater amount or excessive amount of alloying elements and have problems in not only r-value and BH property but also raw material cost.
• the technique for increasing the strength by utilizing microstructure enhancement has problems in production that the prolonged annealing is necessary, and another heat treatment is necessary after the annealing for forming the desired microstructure, and the high-speed cooling equipment is necessary and so on.
• the technique utilizing precipitation of VC or NbC there is still a room for the improvement in the quality stability, productivity, and cost though a high-strength steel sheet having a relatively good workability is obtained.
• the present invention is made with the view of the problems inherent to the conventional techniques and is to provide a high-strength cold rolled steel sheet having excellent deep drawability and bake hardenability, which has not only a tensile strength TS of not less than 440 MPa suitable for use in steel sheets for automobiles and the like but also an average r-value of not less than 1.2 and a bake hardening value (BH value) of not less than 40 MPa as well as an advantageous method for manufacturing the same.
• the high-strength cold rolled steel sheet of the present invention includes a tensile strength of not less than 500 MPa, particularly not less than 590 MPa in addition to the tensile strength of not less than 440 MPa.
• the present invention is a high-strength cold rolled steel sheet having a chemical composition comprising C: 0.010-0.06 mass %, Si: more than 0.5 mass % but not more than 1.5 mass %, Mn: 1.0-3.0 mass %, P: 0.005-0.1 mass %, S: not more than 0.01 mass %, sol. Al: 0.005-0.5 mass %, N: not more than 0.01 mass %, Nb: 0.010-0.090 mass % and Ti: 0.015-0.15 mass % and the remainder being Fe and inevitable impurities, provided that C, Nb, Ti, N and S satisfy the following equations (1) and (2):
• C* C ⁇ (12/93)Nb ⁇ (12/48) ⁇ Ti ⁇ (48/14)N ⁇ (48/32)S ⁇ and each element symbol in the above equations represents a content of each element (mass %), and having a microstructure comprising a ferrite phase with an area ratio of not less than 70% and a martensite phase with an area ratio of not less than 3% and excellent deep drawability and bake hardenability corresponding to a tensile strength of not less than 440 MPa, an average r-value of not less than 1.20 and a BH quantity of not less than 40 MPa.
• the high-strength cold rolled steel sheet of the present invention is characterized by containing not more than 0.5 mass % in total of one or more selected from Mo, Cr and V in addition to the above chemical composition.
• the high-strength cold rolled steel sheet of the present invention is characterized by containing one or two selected from Cu: not more than 0.3 mass % and Ni: not more than 0.3 mass % in addition to the above chemical composition.
• the high-strength cold rolled steel sheet of the present invention is characterized by containing one or two selected from Sn: not more than 0.2 mass % and Sb: not more than 0.2 mass % in addition to the above chemical composition.
• the high-strength cold rolled steel sheet of the present invention is characterized by containing Ta: 0.005-0.1 mass % in addition to the above chemical composition, provided that C, Nb, Ta, Ti, N and S satisfy the following equation (3) instead of the equation (2):
• the present invention proposes a method for manufacturing a high-strength cold rolled steel sheet having excellent deep drawability and bake hardenability by subjecting a steel raw material having a chemical composition comprising C: 0.010-0.06 mass %, Si: more than 0.5 mass % but not more than 1.5 mass %, Mn: 1.0-3.0 mass %, P: 0.005-0.1 mass %, S: not more than 0.01 mass %, sol. Al: 0.005-0.5 mass %, N: not more than 0.01 mass %, Nb: 0.010-0.090 mass % and Ti: 0.015-0.15 mass % and the remainder being Fe and inevitable impurities, provided that C, Nb, Ti, N and S satisfy the following equations (1) and (2):
• C* C ⁇ (12/93)Nb ⁇ (12/48) ⁇ Ti ⁇ (48/14)N ⁇ (48/32)S ⁇ and each element symbol in the above equations represents a content of each element (mass %), to hot rolling, cold rolling and annealing, characterized in that the said annealing is carried out by heating to an annealing temperature of 800-900° C. while a temperature region of 700-800° C. is an average heating rate of less than 3° C./s and thereafter cooling at an average cooling rate of not less than 5° C./s from said annealing temperature to a cooling stop temperature Tc of not higher than 500° C.
• the steel raw material in the manufacturing method of the present invention is characterized by containing not more than 0.5 mass % in total of one or more selected from Mo, Cr and V in addition to the above chemical composition.
• the steel raw material in the manufacturing method of the present invention is characterized by containing one or two selected from Cu: not more than 0.3 mass % and Ni: not more than 0.3 mass % in addition to the above chemical composition.
• the steel raw material in the manufacturing method of the present invention is characterized by containing one or two selected from Sn: not more than 0.2 mass % and Sb: not more than 0.2 mass % in addition to the above chemical composition.
• the steel raw material in the manufacturing method of the present invention is characterized by containing Ta: 0.005-0.1 mass % in addition to the above chemical composition, provided that C, Nb, Ta, Ti, N and S satisfy the following equation (3) instead of the equation (2):
• the manufacturing method of the present invention is characterized in that a rolling reduction of a final pass in a finish rolling of the hot rolling is not less than 10% and a rolling reduction of a pass before the final pass is not less than 15%.
• the manufacturing method of the present invention is characterized in that cooling is started within 3 seconds after the finish rolling of the hot rolling and carried out up to a temperature zone of not higher than 720° C. at an average cooling rate of not less than 40° C./s and coiling is conducted at a temperature of 500-700° C. and thereafter the cold rolling is carried out at a rolling reduction of not less than 50%.
• a high-strength cold rolled steel sheet having excellent deep drawability and bake hardenability with a tensile strength TS of not less than 440 MPa, an average r-value of not less than 1.20 and a BH value of not less than 40 MPa by limiting C content to 0.010-0.06 mass % and restricting a relation of (Nb/93)/(C/12) between Nb addition amount and C content to less than 0.20 so as to render the reduction of solute C badly exerting on the deep drawability as attained in the conventional extremely-low carbon IF steel to a certain level and further controlling solute C (C*) amount not fixed by Nb and Ti to a given range.
• a high-strength cold rolled steel sheet having excellent deep drawability and bake hardenability with an average r-value of not less than 1.20 and a BH value of not less than 40 MPa by reducing expensive Nb as far as possible and positively utilizing Ti even in high-strength steel sheets having not only a tensile strength of not less than 440 MPa but also not less than 500 MPa, particularly not less than 590 MPa.
• the high-strength cold rolled steel sheet of the present invention when applied to automobile parts, it is possible to increase the strength of the part, which has been difficult to conduct press forming, so that the invention largely contributes to improve collision safety of the automobile body and reduce the weight thereof.
• C content is controlled to a range of C: 0.010-0.06 mass %, which is lower than that of the DP steel sheet using the conventional low-carbon steel as a raw material and higher than that of the conventional extremely low-carbon steel sheet, and further adequate amounts of Nb and Ti are added together with the above C content to ensure the adequate solute C amount, whereby not only the development of ⁇ 111 ⁇ recrystallization texture in the annealing is promoted to increase r-value, but also a proper amount of martensite is formed in the cooling after the annealing to increase the strength and further a high bake hardening value (BH value) can be ensured even after the annealing.
• BH value high bake hardening value
• Nb is effective to finely divide the microstructure of the hot rolled steel sheet because it has an effect of delaying the recrystallization. Further, Nb has a high carbide forming ability and precipitates as NbC in steel at a coiling stage after hot rolling, so that the solute C amount can be reduced before cold rolling and before recrystallization annealing.
• Nb is an expensive element and is also an element. deteriorating the productivity (e.g. rolling property). In the present invention, therefore, the amount of Nb added is restricted to a minimum amount required for finely dividing the texture of the hot rolled steel sheet, while Ti having a high carbide forming ability similar to Nb is utilized for reducing the solute C.
• Nb is added so as to satisfy (Nb/93)/(C/12): less than 0.20 in relation to the C content, and further the solute C amount (C*) not fixed by Nb and Ti is controlled to a range of 0.005-0.025 mass %.
• solute C is said to obstruct. the development of ⁇ 111 ⁇ recrystallization texture.
• the solute C required for the formation of martensite is retained without fixing all C as NbC or TiC, and high r-value is attained.
• the reason of providing such an effect. is not clear at present time, it is considered that when the solute C amount is within the above range, positive effect of precipitating fine NbC and TiC into matrix and storing strain in the vicinity of these precipitates during cold rolling to promote formation of ⁇ 111 ⁇ recrystallized grains in addition to the effect of finely dividing the hot rolled steel sheet becomes larger than negative effect. of affecting solute C on the formation of ⁇ 111 ⁇ recrystallization texture.
• the present invention is a feature that the chemical component. of steel is regulated to an adequate range to control the solute C amount (C*) to a range of 0.005-0.025 mass %, and hence high r-value, high BH and high strength based on dual phase are attained. Also, the present invention is another feature that (Nb/93)/(C/12) is regulated to less than 0.20 and Ti is positively utilized as an alternative, whereby the addition amount of expensive Nb increasing burden of hot rolling or cold rolling is considerably decreased and hence it is possible to industrially and stably manufacture high-strength cold rolled steel sheets having high r-value and high BH property without bringing about the increase of raw material cost and the lowering of the productivity.
• the present invention is made by conducting, further examinations on the above new discoveries.
• the C is an important element required for solid-solution strengthening steel and promoting the formation of dual phase comprising ferrite as a primary phase and martensite as a secondary phase and attaining high strength.
• the C content is less than 0.010 mass %, it is difficult to ensure the sufficient amount of martensite and the tensile strength of not less than 440 MPa aiming at the present invention is not obtained.
• the C content exceeds 0.06 mass %, the amount of the resulting martensite increases and the desired average r-value (not less than 1.20) is not obtained.
• the C content is a range of 0.010-0.06 mass %.
• it is a range of 0.020-0.045 mass %.
• Si more than 0.5 mass % but not more than 1.5 mass %
• Si promotes ferrite transformation, enhances C content in non-transformed austenite and easily forms a dual phase of ferrite and martensite, and is also an element having an excellent solid-solution strengthening property.
• more than 0.5 mass % of Si is added in order to ensure tensile strength of not less than 440 MPa. While the amount of Si added exceeds 1.5 mass %, Si-based oxide is formed on the surface of the steel sheet, which deteriorates phosphatability and coating adhesion of a steel sheet product and corrosion resistance after coating. In the present invention, therefore, Si is more than 0.5 mass % but not more than 1.5 mass %.
• the Si content is preferable to be more than 0.8 mass % for tensile strength of not less than 500 MPa. Further, the Si content is preferable to be not less than 1.0 mass % for tensile strength of not less than 590 MPa.
• Mn is an element improving the hardenability of steel and promoting the formation of martensite, so that it is an element effective for the purpose of increasing the strength.
• Mn content is less than 1.0 mass %, it is difficult to form the desired amount of martensite and there is a fear that the tensile strength of not less than 440 MPa cannot be ensured.
• the Mn content is a range of 1.0-3.0 mass %.
• Mn is preferable to be added in an amount of not less than 1.2 mass % for tensile strength of not less than 500 MPa or not less than 1.5 mass % for tensile strength of not less than 590 MPa.
• P is high in the solid solution strengthening property and is an element effective for increasing the strength of steel.
• the P content is less than 0.005 mass %, the effect is not sufficient and it is rather required to remove phosphorus at the steel-making step and hence the rise of production cost is caused.
• P content exceeds 0.1 mass %, P segregates into grain boundaries and the resistance to secondary working brittleness are deteriorated.
• the C amount segregating into grain boundary for contributing to the increase of BH value is lowered, and there is a fear that the desired BH value can not.
• the P content is a range of 0.005-0.1 mass %.
• P is preferably not more than 0.08 mass %, more preferably not more than 0.05 mass % in view of surely ensuing the BH value.
• an upper limit of S is 0.01 mass %. Preferably, it is not more than 0.008 mass %.
• Al is an element added as a deoxidizer, but effectively acts for increasing the strength because it has a solid solution strengthening.
• a content of Al as sol. Al is less than 0.005 mass %, the above effect is not obtained.
• the content of Al as sol. Al exceeds 0.5 mass %, the rise of raw material cost is caused and surface defect of the steel sheet is also caused. Therefore, the content of Al as sol. Al is a range of 0.005-0.5 mass %. Preferably, it is 0.005-0.1 mass %.
• N content exceeds 0.01 mass %, excessive amount of nitride is formed in steel, whereby not only the ductility and toughness but also the surface properties of the steel sheet are deteriorated. Therefore, the N content is not more than 0.01 mass %.
• Nb finely divides the microstructure of the hot rolled steel sheet, and has an action of precipitating as NbC into the hot rolled steel sheet and fixing a part of solute C existing in steel, and contributes to increase the r-value by such an action, so that it is a very important element in the present invention.
• the finely dividing of the microstructure in the hot rolled steel sheet by Nb addition finely divides the microstructure of the steel sheet after the cold rolling and annealing and increases grain boundary area, so that there is an effect of increasing the amount of C segregated into grain boundaries and enhancing BH value. In order to obtain such effects, it is required to add Nb of not less than 0.010 mass %.
• the amount of Nb added is a range of 0.010-0.090 mass %. It is preferably 0.010-0.075 mass %, more preferably 0.010-0.05 mass %.
• Ti is an important element, in the present invention because it contributes to increase the r-value by fixing C and precipitating into the hot rolled steel sheet as TiC likewise Nb. Also, Ti has an action finely dividing the microstructure of the hot rolled steel sheet, which is smaller than that of Nb, so that the amount of C segregated into grain boundaries is increased through the finely dividing of the microstructure of the steel sheet after the annealing and the increase of grain boundary area, and hence there is an effect of enhancing the BH value. In order to obtain such effects, Ti is necessary to be added in an amount of not less than 0.015 mass %. While, excessive addition exceeding 0.15 mass % brings about the rise of raw material cost and further increases the deformation resistance in the cold rolling and hence the stable production is difficult. Also, the excessive Ti addition decreases the solute C and obstructs the formation of martensite in the cooling step after the annealing likewise Nb. Therefore, the amount of Ti added is a range of 0.015-0.15 mass %.
• the high-strength cold rolled steel sheet according to the present invention is necessary not only to satisfy the above chemical composition but also to contain C, Nb, Ti, N and S so as to satisfy the following equations (1) and (2):
• Nb is an expensive element as compared with Ti and is a cause of increasing the rolling load in the hot rolling and obstructing the production stability.
• martensite is formed in the cooling step after the annealing, so that as mentioned later, it is necessary to keep a given amount of solute C (C*) not fixed by Nb or Ti.
• (Nb/93)/(C/12) and C* are necessary to be controlled to proper ranges from a view point of raw material cost, production stability, steel sheet microstructure and properties of steel sheet. Accordingly, the equations (1) and (2) defining the (Nb/93)/(C/12) and C* are most important indications in the present invention.
• (Nb/93)/(C/12) is an atomic ratio of Nb to C.
• this value is not less than 0.20, the amount of NbC precipitated increases and the load in the hot rolling increases and further the addition amount of expensive Nb becomes larger, which makes disadvantageous in the raw material cost. Therefore, (Nb/93)/(C/12) is less than 0.20.
• C* means the amount of solute C not fixed by Nb and Ti. When this value is less than 0.005 mass %, the given amount of martensite cannot be ensured and it is difficult to attain the tensile strength of not less than 440 MPa. While, when C* exceeds 0.025 mass %, the formation of ⁇ 111 ⁇ recrystallization texture in ferrite phase effective for increasing the r-value is obstructed and god deep drawability is not obtained and further there is caused a fear that the desired BH value is not obtained associated with the increase of martensite phase. Therefore, C* is a range of 0.005-0.025 mass %. Moreover, C* is preferably not more than 0.020 mass % for BH value of not less than 50 MPa, while C* is not more than 0.015 mass % for BH value of not less than 60 MPa.
• the high-strength cold rolled steel sheet according to the present invention can be added with one or more selected from Mo, Cr and V and/or one or two selected from Cu and Ni depending upon the required properties.
• Mo, Cr and V are expensive elements, but same as Mn, they are elements improving the hardenability and also elements effective for stably producing martensite. Such effects develop remarkably when the total amount of the above elements added is not less than 0.1 mass %, so that the addition of not less than 0.1 mass % is preferable. While, when the total amount of Mo, Cr and V added exceeds 0.5 mass %, the above effects are saturated and the rise of raw material cost is caused. Therefore, when these elements are added, the total amount is preferable to be not more than 0.5 mass %.
• Cu is a harmful element that causes breakage in the hot rolling and brings about the occurrence of surface flaw.
• the bad influence of Cu upon the properties of the steel sheet is small, so that the content of not more than 0.3 mass % is acceptable.
• Ni is small in the influence upon the properties of the steel sheet likewise Cu and has an effect of preventing the occurrence of surface flaw through the addition of Cu. This effect can be developed by adding in an amount corresponding to not less than 1 ⁇ 2 of the Cu content. However, when the addition amount of Ni becomes excessive, the occurrence of another surface defect resulted from non-uniform formation of scale is promoted, so that the upper limit of Ni addition amount is preferable to be 0.3 mass %.
• the high-strength cold rolled steel sheet according to the present invention can be added with one or two selected from Sn and Sb and/or Ta.
• Sn not more than 0.2 mass %
• Sb not more than 0.2 mass %
• Sn and Sb can be added for suppressing the nitriding or oxidation of the steel sheet surface or decarbonization of several 10 ⁇ m region from the steel sheet surface produced by oxidation. By suppressing such nitriding, oxidation and decarbonization is controlled the reduction of martensite amount formed on the steel sheet surface and is improved the fatigue properties and surface quality.
• Sn or/and Sb are preferable to be added in an amount of not more than 0.005 mass %, respectively. However, the addition exceeding 0.2 mass % fears the deterioration of toughness, so that if added, it is preferable that the upper limit is 0.2 mass %.
• Ta 0.005-0.1 mass %
• Ta has an action of precipitating as TaC in the hot rolled steel sheet and fixing C likewise Nb and Ti, so that it is an element contributing to increase the r-value.
• it is preferable to be added in an amount of not less than 0.005 mass %.
• the addition exceeding 0.1 mass % increases not only the raw material cost, but also obstructs the formation of martensite in the cooling step after the annealing likewise Nb and Ti, or TaC precipitated in the hot rolled steel sheet enhances the deformation resistance in the cold rolling and deteriorates the productivity. Therefore, if added, the Ta amount is preferable to be a range of 0.005-0.1 mass %.
• C* in the equation (3) is less than 0.005, the given amount of martensite cannot be ensured, and it is difficult to obtain the tensile strength of not less than 440 MPa. While, when C* exceeds 0.025, the formation of ⁇ 111 ⁇ recrystallization texture in ferrite phase effective for increasing the r-value is obstructed and the good deep drawability is not obtained and further there is a fear that the desired BH value cannot be ensured associated with the increase of martensite phase. Moreover, C* is preferably not more than 0.020 for the BH value of not less than 50 MPa, while C* is preferably not more than 0.015 for the BH value of not less than 60 MPa.
• the remainder other than the above components comprises Fe and inevitable impurities.
• the other components may be included within the range not damaging the effects of the present invention. Since oxygen (O) forms a non-metal inclusion and affects badly the quality of the steel sheet, the content is preferable to be reduced to not more than 0.003 mass %.
• microstructure of the high-strength cold rolled steel sheet according to the present invention will be described below.
• the high-strength cold rolled steel sheet of the present invention is required to have a microstructure comprising ferrite phase of not less than 70% as an area ratio and martensite phase of not less than 3% as an area ratio with respect to the whole of the microstructure for satisfying the strength of steel sheet, press formability (particularly deep drawability) and bake hardenability together.
• the high-strength cold rolled steel sheet of the present invention may include pearlite, bainite, retained austenite, carbide and so on as a remaining microstructure other than the ferrite phase and martensite phase, but they are acceptable when the total area ratio is not more than 5%.
• the ferrite phase is a soft phase required for ensuring the press formability of the steel sheet, particularly deep drawability.
• the increase of r-value is attained by developing the ⁇ 111 ⁇ recrystallization texture of the ferrite phase.
• the area ratio of the ferrite phase is less than 70%, it is difficult to provide the average r-value of not less than 1.20 and the good deep drawability cannot be obtained.
• the bake hardenability is interrelated with the amount of solute C in ferrite.
• the area ratio of the ferrite phase is less than 70%, it is difficult to attain the BH value of not less than 40 MPa. Therefore, the area ratio of the ferrite phase is not less than 70%.
• the area ratio of the ferrite phase is preferable to be not less than 80% for more enhancing the average r-value and BH value.
• the area ratio of the ferrite phase exceeds 97%, the strength of the steel sheet lowers and it is difficult to ensure the tensile strength of not less than 440 MPa.
• the term “ferrite” used in the present invention includes bainitic ferrite with a high dislocation density transformed from austenite in addition to a polygonal ferrite.
• the martensite phase is a hard phase required for ensuring the strength of the cold rolled steel sheet according to the present invention.
• the area ratio of martensite phase is less than 3%, the strength of the steel sheet lowers and it is difficult to ensure the tensile strength of not less than 440 MPa, so that the area ratio of the martensite phase is made to not less than 3%.
• the martensite phase is preferable to be not less than 5% as an area ratio for providing the tensile strength of not less than 500 MPa or not less than 590 MPa.
• the area ratio of the martensite phase is not more than 30%, preferably not more than 20%.
• the high-strength cold rolled steel sheet according to the present invention is manufactured by sequentially going through a steel-making step of melting a steel having the above adjusted chemical composition in a converter or the like and shaping into a steel raw material (steel slab) through continuous casting or the like, a hot rolling step of subjecting the steel slab to hot rolling comprising rough rolling and finish rolling to form a hot rolled steel sheet, a cold rolling step of subjecting the hot rolled steel sheet to cold rolling to form a cold rolled steel sheet, and an annealing step of annealing the cold rolled steel sheet to provide predetermined strength, deep drawability and bake hardenability.
• the steel melting process is not particularly limited and can adopt such a well-known melting process that molten steel obtained, for example, in a converter, an electric furnace or the like is subjected to secondary refining such as vacuum degassing treatment or the like to provide a given chemical composition.
• a method of forming a steel slab from the molten steel is preferable the use of a continuous casting method from viewpoint of problem such as segregation or the like, but the steel slab may be formed by an ingot-forming—blooming method, a thin slab casting method or the like.
• the thus obtained steel slab is preferable to be reheated and hot rolled.
• the reheating temperature of the steel slab is preferable to be low from a viewpoint. that ⁇ 111 ⁇ recrystallization texture is developed by coarsening precipitates such as TiC and the like for improving the deep drawability.
• the heating temperature of the slab is preferable to be not lower than 1000° C.
• the upper limit of the heating temperature is preferable to about 1300° C. from a viewpoint of suppressing the increase of scale loss due to oxidation.
• the slab In the hot rolling of the steel slab, it is common that the slab is charged into a heating furnace and reheated to a given temperature and then rolled. However, when the slab after the continuous casting is above the given temperature, the slab may be rolled (direct rolling) without reheating as it is, or there may be adopted a method wherein the slab is charged into a heating, furnace at a higher temperature state and a part of the reheating is omitted (warm piece charging).
• the steel slab reheated under the above conditions is subjected to rough rolling to form a sheet bar.
• the rough rolling conditions are not particularly defined because it may be conducted according to the usual manner.
• the temperature of the sheet bar may be increased by utilizing a sheet bar heater in view of ensuring a given hot rolling temperature or preventing troubles of rolling.
• the sheet bar after the rough rolling is then subjected to finish rolling to form a hot rolled steel sheet.
• the fine division of the microstructure in the hot rolled steel sheet increases preferential nucleation sites of ⁇ 111 ⁇ recrystallization texture in the annealing after cold rolling, so that it is effective for improving the r-value.
• the rolling reduction of the final pass is less than 10%, ferrite grains are coarsened and there is a fear that the effect of increasing the r-value or BH value is not obtained. Therefore, the rolling reduction of the final pass is preferably not less than 10%, more preferably not less than 13%.
• the rolling reduction of a pass before the final pass is not less than 15% in addition to the aforementioned control of the rolling reduction of the final pass.
• the rolling reduction of the pass before the final pass is more enhanced strain accumulation effect, whereby many shear bands are introduced into old austenite grains and hence nucleation sites of ferrite transformation are further increased and the microstructure of the hot rolled steel sheet is finely divided to further improve the r-value and BH value.
• the rolling reduction of the pass before the final pass is preferably not less than 15%, more preferably not less than 18%.
• the upper limit of the rolling reduction in the final pass and the pass before the final pass is preferable to be less than 40% in view of the rolling load.
• the rolling temperatures of the final pass and the pass before the final pass are not particularly limited.
• the rolling temperature of the final pass is preferably not lower than 800° C., more preferably not lower than 830° C.
• the rolling temperature of the pass before the final pass is preferably not higher than 980° C., more preferably not higher than 950° C.
• the transformation from non-recrystallized austenite to ferrite becomes larger and the microstructure of the steel sheet after the cold rolling and annealing is influenced by the microstructure of the hot rolled steel sheet and forms a non-uniform microstructure elongated in a rolling direction and the workability is deteriorated.
• the strain accumulation effect becomes insufficient due to recovering and it is difficult to finely divide the microstructure of the hot rolled steel sheet, and the effect of increasing the r-value or BH value may not be obtained.
• the hot rolled steel sheet after the hot rolling is started to cooling within 3 seconds after the finish rolling and cooled at an average cooling rate of not less than 40° C./s to a temperature region of not higher than 720° and then coiled at a temperature of 500-700° C. in view of the improvement of r-value or BH value by fine division of crystal grains.
• the microstructure of the hot rolled steel sheet becomes coarse, and the effect of increasing the r-value or BH value may not be obtained.
• the coiling temperature exceeds 700° C.
• the microstructure of the hot rolled steel sheet is coarsened, and there is a fear of lowering the strength and the increase of the r-value or BH value may be obstructed after the cold rolling and annealing.
• the coiling temperature is lower than 500° C.
• the precipitation of NbC or TiC is difficult and the solute C is increased, which is disadvantageous in the increase of the r-value.
• the hot rolled steel sheet is then pickled and cold-rolled according to the usual manner to form a cold rolled steel sheet.
• the rolling reduction in the cold rolling is preferable to be a range of 50-90%.
• the rolling reduction in the cold rolling is more preferable to be set to a higher level.
• the rolling reduction is less than 50%, the ⁇ 111 ⁇ recrystallization texture of the ferrite phase is not developed sufficiently and the excellent. deep drawability may not be obtained.
• the rolling reduction exceeds 90% the load in the cold rolling is increased and there is a fear of causing troubles in the passing of the sheet.
• the cold rolled steel sheet is then annealed to provide desirable strength, deep drawability and bake hardenability.
• the steel sheet is heated to an annealing temperature of 800-900° C. at an average heating rate of less than 3° C./s within a temperature range of 700-800° C., soaked and then cooled from the annealing temperature (soaking temperature) to a cooling stop temperature Tc of not higher than 500° C. at an average cooling rate of not less than 5° C./s.
• the annealing method satisfying the above conditions is preferably adapted continuous annealing.
• the recrystallization temperature of the steel sheet after the cold rolling becomes relatively high.
• the average heating rate is not less than 3° C./s, the development of ⁇ 111 ⁇ recrystallization texture is insufficient and the increase of the r-value may be difficult.
• the average heating rate is preferable to be not less than 0.5° C./s in view of the enhancement of productivity.
• the annealing temperature (soaking temperature) is necessary to be a two-phase zone temperature of ferrite phase and austenite phase.
• the annealing temperature is a temperature range of 800-900° C.
• the annealing temperature is lower than 800° C.
• the desired martensite quantity is not obtained after the cooling followed by the annealing, and also the recrystallization is not sufficiently completed during the annealing and hence there is a fear that the ⁇ 111 ⁇ recrystallization texture of the ferrite phase is not developed and the average r-value of not less than 1.20 cannot be obtained.
• the annealing temperature exceeds 900° C., the amount of solute C in ferrite decreases, and there is a fear that BH value of not less than 40 MPa cannot be ensured.
• the annealing temperature exceeds 900° C.
• secondary phase martensite phase, bainite phase, pearlite phase
• the annealing temperature is a range of 800-900° C., preferably a range of 820-880° C.
• the time for keeping the soaking in the annealing is preferable to be not less than 15 seconds (s) in view that enrichment of an element such as C or the like in austenite proceeds sufficiently and that the development of ⁇ 111 ⁇ recrystallization texture of the ferrite phase is promoted sufficiently.
• the time for keeping the soaking exceeds 300 seconds (s)
• the grains are coarsened and the high BH value is not obtained and there is a fear of causing had influence on the properties of the steel sheet such as lowering of strength, deterioration of surface properties of steel sheet and the like. Therefore, the time for keeping the soaking is preferably a range of 15-300 seconds (s), more preferably a range of 15-200 seconds (s).
• the steel sheet after the completion of recrystallization in the annealing is necessary to be cooled from the annealing temperature (soaking temperature) to a cooling stop temperature Tc of not higher than 500° C. at an average cooling rate of not less than 5° C./s.
• Tc cooling stop temperature
• the average cooling rate is less than 5° C./s, it is difficult to ensure martensite phase of not less than 3% as an area ratio with respect. to the whole microstructure of the steel sheet and the desired strength (tensile strength of not less than 440 MPa) may not be obtained.
• the cooling stop temperature exceeds 500° C., the martensite phase of not less than 3% as an area ratio may not still obtained.
• the average cooling rate is preferably not less than 8° C./s, more preferably not less than 10° C./s, and the cooling stop temperature Tc is preferably a range of 400-450° C. If the average cooling rate exceeds 100° C./s, special equipment for water cooling or the like is required, which brings about the increase of production cost and the deterioration of steel sheet form, so that the upper limit of the average cooling rate is preferable to be about 100° C./s.
• the cooling conditions other than the cooling stop temperature Tc are not particularly limited. However, in order to properly proceed the tempering of martensite phase to recover the ductility and toughness, it is preferable to conduct cooling in a temperature region of from the cooling stop temperature Tc to 200° C. at an average cooling rate of 0.2-10° C./s. When the average cooling rate in the above temperature region is less than 2° C./s, the tempering of the martensite phase proceeds excessively and the desired strength may not be obtained. While, when the average cooling rate in the above temperature region exceeds 10° C./s, the tempering of the martensite phase does not proceed sufficiently and the effect of recovering the ductility and toughness cannot be expected. More preferably, the average cooling rate is a range of 0.5-6° C./s.
• the cold rolled steel sheet of the present invention manufactured as mentioned above may be subsequently subjected to temper rolling, leveler work or the like for the purpose of correcting the form, adjusting surface roughness, and so on.
• temper rolling rate is preferable to be about 0.3-1.5%.
• Each of steels A-V having a chemical composition shown in Table 1 is melted by a well-known refining process including converter, vacuum degassing treatment and the like and then continuously cast to form a steel slab of 260 mm in thickness.
• Each of these steel slabs is heated to 1220° C. and hot rolled to obtain a hot rolled steel sheet having a thickness of 3.8 mm.
• the rolling temperature and rolling reduction of each of final pass and pass before final pass in finish rolling during the hot rolling, average cooling rate from cooling start to 720° C. after the completion of the finish rolling and coiling temperature are shown in Table 2, and the time after the completion of the finish rolling to the start of cooling is within 3 seconds.
• the hot rolled steel sheet is pickled and cold rolled under conditions shown in Table 2 to form a cold rolled steel sheet having a thickness of 1.2 mm, which is continuously annealed under conditions shown in Table 2 and subjected to temper rolling at 0.5% to obtain a cold rolled steel sheet (product).
• a sample is taken from each of the thus obtained cold rolled steel sheets and subjected to microstructure observation and tensile test by the following methods, whereby the steel sheet microstructure is identified and the area ratios of ferrite phase and martensite phase, tensile strength, elongation, average r-value and bake hardening value (BH value) are measured.
• BH value bake hardening value
• a test piece for the microstructure observation is taken from the above sample and L-section (vertical section parallel to the rolling direction) is mechanically polished and corroded with nital and shot with a scanning type electron microscope (SEM) at a magnification of 2000 times to obtain a microstructure photograph (SEM photograph), from which the steel sheet microstructure is identified and area ratios of ferrite phase and martensite phase are measured.
• SEM photograph a microstructure photograph
• the identification of the steel sheet microstructure from the above photograph indicates that ferrite is a slight black contrast region, pearlite is a region forming a carbide in lamellar form, bainite is a region forming a carbide in dot sequence, and martensite and retained austenite (retained ⁇ ) are particles with white contrast.
• test piece is subjected to tempering treatment of 250° C. ⁇ 4 hr and shot in the same manner as mentioned above to obtain a microstructure photograph, from which area ratios are again determined when a region forming a carbide in lamellar form is a pearlite region before heat treatment and a region forming a carbide in dot sequence is a bainite or martensite region before heat treatment, and fine particles remaining with white contrast are measured as retained ⁇ , and the area ratio of martensite phase is determined by subtracting from area ratio of particles with white contrast before tempering treatment (martensite and retained austenite).
• the area ratio of each phase is colored separately every each phase on a transparent OHP sheet and incorporated into an image and binarized to measure an area ratio by an image analyzing software (Digital Image-Pro Plus ver. 4.0 made by Microsoft).
• JIS No. 5 tensile specimen (JIS Z2201), wherein tension direction is 90° direction (C-direction) with respect to the rolling direction, is taken from the above sample and subjected to a tensile test according to a definition of JIS 22241 to measure tensile strength TS and total elongation E 1 .
• the bake hardening value (BH value) is obtained by applying a tensile pre-strain of 2%, subjecting to a heat treatment corresponding to coat baking conditions of 170° C. ⁇ 20 minutes, again conducting the tensile test and measuring a value of subtracting nominal stress in the application of the pre-strain from an upper yield point. after the heat treatment as BH value.
• JIS No. 5 tensile specimens wherein tensile directions are 0° direction (L-direction), 45° direction (D-direction) and 90° direction (C-direction) with respect to the rolling direction, are taken from the above sample, and then true strain in widthwise direction and true strain in thickness direction of each specimen are measured when uniaxial tensile strain of 10% is applied to each of these specimens, and an average r-value is calculated from these measured values according to the definition of JIS Z2254.
• Steel sheets of Nos. 3-13 and 16-22 are Invention Examples wherein the chemical composition and production conditions are acceptable in the present invention, and have a tensile strength TS of not less than 440 MPa, an average r-value of not less than 1.20 and a BH value of not less than 40 MPa, so that they are cold rolled steel sheets satisfying the strength, deep drawability and bake hardening value.
• the steel sheets of Nos. 8, 12, 13 and 22 having a solute C amount (C*) of not more than 0.020 mass % have BH value of not less than 50 MPa
• the steel sheets of Nos. 3-7 and 16-20 having C* of not more than 0.015% have BH value of not less than 60 MPa, so that they have a very high bake hardening value.
• the steel slab having a chemical composition for steels D, G and L shown in Table 1 is heated to 1220° C. and then hot rolled to form a hot rolled steel sheet of 3.8 mm in thickness.
• the finish rolling conditions, cooling conditions and coiling temperature in the hot rolling are shown in Table 4.
• the time of from the completion of the finish rolling to the cooling start is within 3 seconds.
• the hot rolled steel sheet is pickled and cold rolled under conditions shown in Table 4 to form a cold rolled steel sheet of 1.2 mm in thickness, which is continuously annealed under conditions shown in Table 4 and subjected to temper rolling at 0.5% to obtain a cold rolled steel sheet product.
• a test piece is taken from the thus obtained cold rolled steel sheet in the same manner as in Example 1 and subjected to microstructure observation and tensile test, and also area ratios of ferrite, martensite and the like, tensile strength, elongation, average r-value and bake hardening value are measured.
• the steel sheets of Invention Example Nos. 23-29, 31, 32, 35, 36, 38 and 39 satisfying the production conditions of the present invention are steel sheets having a tensile strength TS of not less than 440 MPa, an average r-value of not less than 1.20 and a BH value of not less than 40 MPa and satisfying the strength, deep drawability and bake hardenability.
• the high-strength cold rolled steel sheet of the present invention is not limited to application for automobile members and can be preferably used in other applications requiring high strength, deep drawability and bake hardenability. Therefore, it is suitable as a material for household electrical goods, steel pipes and so on.

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)

Applications Claiming Priority (5)

Publications (1)

Family

ID=46024408

Family Applications (1)

Country Status (10)

Cited By (7)

Families Citing this family (6)

Citations (3)

Family Cites Families (18)

• 2011
  • 2011-10-17 JP JP2011227817A patent/JP5825481B2/ja active Active
  • 2011-10-28 MX MX2013005011A patent/MX350226B/es active IP Right Grant
  • 2011-10-28 WO PCT/JP2011/074939 patent/WO2012060294A1/ja active Application Filing
  • 2011-10-28 CA CA2814193A patent/CA2814193C/en not_active Expired - Fee Related
  • 2011-10-28 EP EP11837944.5A patent/EP2636762B1/en active Active
  • 2011-10-28 US US13/882,300 patent/US20130213529A1/en not_active Abandoned
  • 2011-10-28 KR KR1020137009954A patent/KR101561358B1/ko not_active IP Right Cessation
  • 2011-10-28 CN CN201180053425.3A patent/CN103201403B/zh active Active
  • 2011-10-28 BR BR112013011013A patent/BR112013011013A2/pt not_active Application Discontinuation
  • 2011-11-03 TW TW100140129A patent/TWI473887B/zh not_active IP Right Cessation

Patent Citations (3)

Also Published As

Similar Documents

Legal Events