WO2018163871A1 - 高強度熱延めっき鋼板 - Google Patents
高強度熱延めっき鋼板 Download PDFInfo
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- WO2018163871A1 WO2018163871A1 PCT/JP2018/006847 JP2018006847W WO2018163871A1 WO 2018163871 A1 WO2018163871 A1 WO 2018163871A1 JP 2018006847 W JP2018006847 W JP 2018006847W WO 2018163871 A1 WO2018163871 A1 WO 2018163871A1
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- C—CHEMISTRY; METALLURGY
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- 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
- C21D8/0205—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips of ferrous alloys
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- B—PERFORMING OPERATIONS; TRANSPORTING
- B32—LAYERED PRODUCTS
- B32B—LAYERED PRODUCTS, i.e. PRODUCTS BUILT-UP OF STRATA OF FLAT OR NON-FLAT, e.g. CELLULAR OR HONEYCOMB, FORM
- B32B15/00—Layered products comprising a layer of metal
- B32B15/01—Layered products comprising a layer of metal all layers being exclusively metallic
- B32B15/013—Layered products comprising a layer of metal all layers being exclusively metallic one layer being formed of an iron alloy or steel, another layer being formed of a metal other than iron or aluminium
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- C21D6/00—Heat treatment of ferrous alloys
- C21D6/002—Heat treatment of ferrous alloys containing Cr
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- C21D6/005—Heat treatment of ferrous alloys containing Mn
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- 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
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- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
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- 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
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- C22C38/001—Ferrous alloys, e.g. steel alloys containing N
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/002—Ferrous alloys, e.g. steel alloys containing In, Mg, or other elements not provided for in one single group C22C38/001 - C22C38/60
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/005—Ferrous alloys, e.g. steel alloys containing rare earths, i.e. Sc, Y, Lanthanides
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
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- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
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- C22C38/06—Ferrous alloys, e.g. steel alloys containing aluminium
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- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
- C22C38/14—Ferrous alloys, e.g. steel alloys containing titanium or zirconium
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- 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
- C22C38/22—Ferrous alloys, e.g. steel alloys containing chromium with molybdenum or tungsten
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- 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
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- 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
- C22C38/32—Ferrous alloys, e.g. steel alloys containing chromium with boron
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- 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
- C22C38/38—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of manganese
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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
- C23C2/06—Zinc or cadmium or alloys based thereon
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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/26—After-treatment
- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
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- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
- C23C2/36—Elongated material
- C23C2/40—Plates; Strips
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/002—Bainite
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/004—Dispersions; Precipitations
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- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
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- Y—GENERAL TAGGING OF NEW TECHNOLOGICAL DEVELOPMENTS; GENERAL TAGGING OF CROSS-SECTIONAL TECHNOLOGIES SPANNING OVER SEVERAL SECTIONS OF THE IPC; TECHNICAL SUBJECTS COVERED BY FORMER USPC CROSS-REFERENCE ART COLLECTIONS [XRACs] AND DIGESTS
- Y10—TECHNICAL SUBJECTS COVERED BY FORMER USPC
- Y10T—TECHNICAL SUBJECTS COVERED BY FORMER US CLASSIFICATION
- Y10T428/00—Stock material or miscellaneous articles
- Y10T428/12—All metal or with adjacent metals
- Y10T428/12493—Composite; i.e., plural, adjacent, spatially distinct metal components [e.g., layers, joint, etc.]
- Y10T428/12771—Transition metal-base component
- Y10T428/12785—Group IIB metal-base component
- Y10T428/12792—Zn-base component
- Y10T428/12799—Next to Fe-base component [e.g., galvanized]
Definitions
- the present invention relates to a high-strength hot-rolled plated steel sheet excellent in formability and suitable for use as a material for structural members such as automobiles, transportation equipment, and building equipment.
- Patent Document 1 in mass%, C: 0.03% to 0.35%, Si: 0.01% to 2.0%, Mn: 0.3% to 4.0% , P: 0.001% to 0.10%, S: 0.0005% to 0.05%, N: 0.0005% to 0.010%, Al: 0.01% to 2.0% %, With the balance being Fe and inevitable impurities, and by using recrystallization control in hot rolling and reheating performed after runout table cooling or hot rolling, the ferrite grain coverage by the martensite phase is 30% A method for obtaining a high-strength hot-rolled steel sheet having an ultra-thin metal structure and excellent in ductility and stretch flangeability and having a tensile strength of 500 MPa or more is disclosed.
- Patent Document 2 includes mass%, C: 0.02% to 0.075%, Si: 0.001% to 0.2%, Mn: 2.0% to 4.5%, P : 0.1% or less, S: 0.01% or less, sol. Al: 0.001% or more and 0.2% or less, N: 0.01% or less, O: 0.01% or less, and addition amount of Ti and Nb is 0.14 ⁇ Ti + Nb / 2 ⁇ 0.3
- the component composition satisfies the above, and by adding a large amount of Mn compared to the conventional steel plate, the formation of coarse cementite, martensite and austenite phases in the hot dip galvanizing process is suppressed, and the stretch flangeability is excellent.
- a method for obtaining a high-strength hot-dip galvanized steel sheet having a strength of 590 MPa or more is disclosed.
- Patent Document 3 in mass%, C: 0.03% to 0.12%, Si: 0.01% to 0.5%, Mn: 1.4% to 5.0%, P : 0.05% or less, S: 0.010% or less, sol.
- a hot-rolled steel sheet having a chemical composition containing Al: 0.001% or more and 0.5% or less and N: 0.020% or less is heated to a temperature range of 650 ° C. or more and 950 ° C. or less and 3 ° C./second or more and 20 By cooling to 550 ° C.
- Patent Document 4 includes mass%, C: 0.02% to 0.10%, Si: 0.005% to 0.5%, Mn: 1.4% to 2.5%, P : 0.025% or less, S: 0.010% or less, sol. Al: 0.001% or more and 0.2% or less, N: 0.008% or less and Ti: 0.15% or less, Ca: 0.01% or less, Mg: 0.01% or less, and REM And containing one or more selected from the group consisting of 0.01% or less, having a chemical composition satisfying the following formulas (3) to (5), and having an area% of ferrite: 50 to 94%, bainite: 5 to 49% and martensite and retained austenite total: 1 to 20% of steel structure, tensile strength excellent in ductility and stretch flangeability: high strength hot dip galvanizing of 490 MPa or more A method for obtaining a steel sheet is disclosed.
- Patent Document 5 C: 0.025 to 0.15%, Si: 0.01 to 1.0%, Mn: 1.0 to 2.5%, P: 0.02% or less, S: 0 0.005% or less, Al: 0.5% or less, Ti: 0.04 to 0.10%, and N: 0.007% or less, with the balance consisting of Fe and inevitable impurities, Mn / Ti ratio: having a composition of 15 or more, providing a waiting time of 45 seconds or more from rough rolling to finish rolling, and setting the reduction rate of the final three stages of finish rolling to 25% or more, corresponding to crystal grains
- Disclosed is a method for suppressing the occurrence of punching fractures and uneven surfaces in rolled steel
- Patent Document 1 it is necessary to provide intermediate cooling in runout table cooling, and there is a problem that characteristics are not stable over the entire length of the steel sheet. Further, a hard second phase such as martensite is likely to be a starting point for punching end face cracking during punching, but this point has not been studied at all.
- the Mn content is larger than that of the conventional steel plate, Mn oxide is easily generated on the steel plate surface in the annealing process in the plating process, and the steel plate surface is stably plated over the entire length of the steel plate. There is a problem that it is not easy to form a layer.
- Patent Documents 3 and 4 there is a problem that a punched end face is likely to be formed because it has a metal structure containing at least 1% of a hard phase such as martensite and retained austenite.
- a hard phase such as martensite and retained austenite.
- Ca, Mg, and REM are added for the purpose of improving the hole expansion property, but these are coarse if the addition amount balance of O and S contained in the steel sheet is not appropriate. This forms an oxide, which becomes the starting point of cracks and conversely deteriorates hole expandability and punched end face properties.
- the object of the present invention is to solve such problems and to provide a high-strength hot-rolled plated steel sheet excellent in formability.
- the tensile strength is 570 MPa even when a ferrite phase having a low dislocation density or a tempered bainite phase is used as a main phase. It was found that the above high strength can be obtained stably. In addition, it was also found that, by strengthening the ferrite phase and tempered bainite phase with fine metal carbide, the hardness difference between the main phase and the second phase is reduced, and the stretch flangeability is further improved.
- the present invention has been made on the basis of the above findings, and has been made as a result of repeated studies on the optimum component composition and metal structure.
- the gist of the present invention is as follows. [1] By mass%, C: 0.03 to 0.15%, Si: 0.4% or less, Mn: 1.2 to 1.9%, Ti: 0.05 to 0.25%, B: 0.0005 to 0.0050%, P: 0.03% or less, S: 0.005% or less, Al: 0.005 to 0.4%, N: 0.01% or less, with the balance being Fe And a component composition comprising inevitable impurities, and at least one of ferrite and tempered bainite has a total area ratio of 90% or more, and the volume fraction of Ti carbide having a particle diameter of 20 nm or less is 0.05 vol% or more.
- the composition further contains, in mass%, one or more of Nb, V, Zr, Mo, Cr, W, Ta, and Hf in a total amount of 0.1% or less.
- by mass% at least one of Ca: 0.005% or less, Mg: 0.005% or less, REM: 0.005% or less, The high-strength hot-rolled plated steel sheet according to the above [1] or [2], which is contained so as to satisfy (1).
- “high strength” means that the tensile strength (TS) is 570 MPa or more.
- “Excellent formability” means excellent punchability and stretch flangeability, and “Excellent punchability” means that the arithmetic average roughness Ra of the fractured surface portion appearing on the punching end face is 15 ⁇ m or less.
- “Excellent stretch flangeability” means that the critical hole expansion rate ( ⁇ ) (sometimes simply referred to as the hole expansion rate) is 60% or more.
- the “steel plate” includes a steel plate and a steel strip.
- the above-mentioned problems of the prior art are advantageously solved, and a high-strength hot-rolled plated steel sheet excellent in formability can be obtained. It has a high tensile strength of 570 MPa or more and has good punchability, high stretch flangeability, and good plating properties, so it is suitable as a material for structural members such as automobiles, transportation equipment, and building equipment. .
- the present invention is an industrially extremely useful invention.
- C 0.03-0.15%
- C is an element that contributes to strength improvement by forming metal carbide in the steel sheet. If the C content is less than 0.03%, the desired strength cannot be obtained. On the other hand, if the C content exceeds 0.15%, the amount of bainite, pearlite, cementite, martensite, and retained austenite generated in the steel sheet increases, and the punching workability and stretch flangeability deteriorate. For this reason, the C content is 0.03 to 0.15%. Preferably it is 0.05 to 0.13%, more preferably 0.06 to 0.13%, still more preferably 0.07 to 0.11%.
- Si 0.4% or less Si is an effective element as an element for improving the strength without reducing the ductility.
- the adhesion of the plating layer is inhibited and the corrosion resistance of the steel sheet is deteriorated.
- the upper limit of Si content is made 0.4%.
- it is 0.2% or less, More preferably, it is 0.05% or less. Note that there is no problem even if the Si content is zero.
- Mn 1.2 to 1.9%
- Mn is an element that contributes to improving the strength of the steel sheet through solid solution strengthening and grain refinement strengthening. If the Mn content is less than 1.2%, the desired strength cannot be obtained. On the other hand, if the Mn content exceeds 1.9%, the hardenability is excessively increased, coarse martensite and retained austenite are generated, and the punching workability and stretch flangeability are deteriorated. Therefore, the Mn content is 1.2 to 1.9%.
- the content is preferably 1.40 to 1.85%, more preferably 1.41 to 1.80%.
- Ti 0.05 to 0.25%
- Ti is an important element in the present invention that precipitates as fine Ti carbide during the annealing process before the plating process and contributes to the enhancement of the strength of the steel sheet.
- the Ti content is set to 0.05 to 0.25%.
- it is 0.05 to 0.20%, more preferably 0.05 to 0.15%.
- B 0.0005 to 0.0050% B is added to improve hardenability and suppress excessive formation of bainite, pearlite, and cementite.
- the B content needs to be 0.0005% or more.
- the B content is set to 0.0005 to 0.0050%.
- it is 0.0005 to 0.0040%, more preferably 0.0005 to 0.0030%.
- P 0.03% or less
- P is an element contained as an impurity, and segregates at the prior austenite grain boundaries to reduce toughness. Therefore, in order to improve punchability and stretch flangeability, it is preferable to reduce as much as possible. However, the content up to 0.03% is acceptable. For this reason, the P content is set to 0.03% or less. Preferably it is 0.02% or less. Note that there is no problem even if the P content is zero.
- S 0.005% or less S forms Ti sulfide and inhibits the strength improvement effect by addition of Ti. Moreover, Mn sulfide is formed, and punchability and stretch flangeability are deteriorated. For this reason, although it is preferable to reduce S content as much as possible, the content to 0.005% or less is permissible. For this reason, S content shall be 0.005% or less. Preferably, it is 0.004% or less, more preferably 0.003% or less. Note that there is no problem even if the S content is zero.
- Al acts as a deoxidizer and is an effective element for improving the cleanliness of steel.
- the Al content needs to be 0.005% or more.
- the Al content is set to 0.005 to 0.4%.
- the content is 0.005 to 0.1%, more preferably 0.01 to 0.06%.
- N 0.01% or less N is likely to bond with Ti at a high temperature to form coarse nitrides, and deteriorates punching workability and stretch flangeability. For this reason, N content shall be 0.01% or less. Preferably it is 0.008% or less. More preferably, it is 0.006% or less. There is no problem even if the N content is zero.
- the composition may further contain 0.1% or less in total of one or more of Nb, V, Zr, Mo, Cr, W, Ta, and Hf in mass%.
- Nb, V, Zr, Mo, Cr, W, Ta, and Hf 0.1% or less in total of one or more of Nb, V, Zr, Mo, Cr, W, Ta, and Hf in mass%.
- Each of these elements alone or in combination with Ti forms carbides and contributes to increasing the strength. However, if the total content exceeds 0.1%, formation of fine Ti carbides may be hindered and strength may be reduced.
- the mass% is Ca: 0.005% or less, Mg: 0.005% or less, REM: 0.005% or less.
- the mass% is Ca: 0.005% or less, Mg: 0.005% or less, REM: 0.005% or less.
- One or more of them can be contained in total so as to satisfy the following formula (1).
- the formula (1) is preferably 0.8 or more and 4.5 or less, more preferably 1.0 or more and 4.0 or less.
- the balance is Fe and inevitable impurities.
- elements mixed from ore or scrap Cu, Ni, Sb, As, Sn, Pb, etc.
- the steel sheet of the present invention has a plating layer on the surface for the purpose of improving corrosion resistance in order to make it a hot rolled steel sheet suitable as a material for automobile parts exposed to severe corrosive environments.
- the type of the plating layer is not particularly limited, and electroplating or hot dipping may be used. If it is hot dip plating, a hot dip galvanization is mentioned as a suitable example.
- the plating layer may be alloyed plating subjected to alloying treatment.
- the high-strength hot-rolled plated steel sheet of the present invention has at least one of ferrite and tempered bainite as the main phase, and the area ratio of the main phase is 90% or more in total. Further, Ti carbide having a particle size of 20 nm or less is finely dispersed, and the volume fraction of Ti carbide having a particle size of 20 nm or less is 0.05 vol% or more.
- the main phase is a structure having a low dislocation density. That is, in the present invention, the main phase is any of the case where ferrite is the main phase, the case where tempered bainite is the main phase, and the case where ferrite and tempered bainite are the main phase.
- the area ratio of the main phase is 90% or more in total, punchability and stretch flangeability are improved. Preferably it is 92% or more.
- a structure in which the grain orientation spread obtained by the SEM / EBSD method is 4 ° or less is defined as a structure having a low dislocation density.
- This Grain Orientation Spread arithmetically averages the orientation difference between a certain measurement point in one crystal grain and the other measurement points in the same crystal grain, and performs the same calculation for all the measurement points in the same crystal grain. It is a value obtained by arithmetically averaging them.
- the value is correlated with the average strain information in the crystal grains, that is, the dislocation density.
- a hot-rolled steel sheet mainly composed of ferrite and bainite generated by phase transformation in the hot rolling process is reheated to an Ac3 point or less.
- a tempered bainite that is generated when the steel sheet is deformed, or a ferrite phase that is formed by phase transformation from reverse-transformed austenite when the hot-rolled steel sheet is reheated to the Ac3 point or lower.
- examples of the structure contained in the steel sheet include bainite, pearlite, martensite, retained austenite, and the like as structures other than the above ferrite and tempered bainite, and these structures have a total area ratio of 10% or less.
- the area ratio is less than 1% in total, especially for coarse forms with a major axis of 1 ⁇ m or more, the area ratio is The total amount is preferably less than 1%.
- the amount of Ti carbide is also important for fine dispersion strengthening.
- the volume fraction of Ti carbide having a particle diameter of 20 nm or less is less than 0.05 vol%, a desired strengthening amount cannot be obtained.
- a MC is the atomic weight of Ti carbide (g / mol)
- AM is the atomic weight of Ti (g / mol)
- ⁇ Fe is the density of ⁇ iron (g / cm 3 )
- ⁇ MC is Ti
- the density (g / cm 3 ) of carbide and [% M P ] represent the mass concentration with respect to all added elements of Ti present in Ti carbide having a particle diameter of 20 nm or less. [% M P ] is determined by inductively coupled plasma atomic emission spectrometry after fractionating 20 nm or less of Ti carbide by the electrolytic extraction residue method.
- Nb, V, Zr, Mo, Cr, W, Ta, Hf may be contained in the Ti carbide. Even when Ti carbide contains other elements or metal carbides other than Ti carbide, the physical properties and mass concentration used to calculate the volume fraction f in formula (2) use Ti and Ti carbide values. To do. In addition to Ti carbide, a single metal carbide composed of Nb, V, Zr, Mo, Cr, W, Ta, and Hf may exist. In the present invention, carbides of Ti and metals such as Nb, V, Zr, Mo, Cr, W, Ta, and Hf are also Ti carbides. In the present invention, the volume fraction of these carbides is not incorporated in the metal structure fraction.
- the high-strength hot-rolled plated steel sheet of the present invention is subjected to hot rolling consisting of rough rolling and finish rolling to the steel material having the above-described component composition, and after finishing rolling is cooled, wound, and hot-rolled steel sheet, Subsequently, it can manufacture by annealing a hot-rolled steel plate and making a plating layer adhere.
- the method for melting the steel material is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed.
- a slab steel material
- a continuous casting method because of problems such as segregation, but as a slab by a known casting method such as an ingot-bundling rolling method or a thin slab continuous casting method. Also good.
- hot rolling the slab after casting it may be rolled after reheating the slab in a heating furnace, or when the slab is maintained at a temperature higher than a predetermined temperature, it is not reheated. Direct rolling may be used.
- heating temperature when heating a slab, heating temperature shall be 1150 degreeC or more. However, when the heating temperature becomes excessively high, the surface is excessively oxidized and TiO 2 is generated, so that Ti is consumed, and in the case of a steel plate, the strength of the surface layer is likely to be lowered. 1350 ° C. or less.
- the slab before rough rolling maintains a temperature equal to or higher than a predetermined temperature and the carbide in the slab is sufficiently dissolved, the step of heating the steel material before rough rolling is omitted. can do.
- the rough rolling conditions are not particularly limited.
- the present invention in order to obtain a high-strength hot-rolled plated steel sheet having a metal structure in which fine Ti carbide particles having a particle size of 20 nm or less are dispersed, it is important to suppress Ti carbide precipitation in the hot rolling process as much as possible. is there. If a large amount of Ti carbide precipitates in the hot rolling process, the growth and coarsening of the Ti carbide proceed in the subsequent annealing process, and the desired strength cannot be obtained. Further, a hardness difference between the main phase and the second phase occurs, and the punching workability and stretch flangeability deteriorate.
- finish rolling is performed in a temperature range of 840 ° C. or higher.
- the finish rolling temperature is less than 840 ° C.
- the ferrite transformation is likely to proceed during rolling, and metal carbides such as Ti carbide precipitate at the same time as the ferrite transformation.
- cooling after finish rolling cooling is performed so that the average cooling rate between 780 ° C. and 680 ° C. is 30 ° C./s or more. When the average cooling rate is less than 30 ° C./s, ferrite transformation proceeds during cooling, and Ti carbide precipitates.
- the winding temperature is 350 ° C. or higher and 550 ° C. or lower.
- the coiling temperature is less than 350 ° C.
- the metal structure becomes a martensite phase having an extremely high dislocation density, and 90% or more of a structure having a low dislocation density is obtained in which the grain orientation spread is 4 ° or less even after the subsequent annealing step. It becomes difficult.
- the coiling temperature exceeds 550 ° C., ferrite transformation proceeds in the coiling process, Ti carbide precipitates in the hot-rolled steel sheet, and Ti carbide grows in the annealing process when depositing the plating layer. Strength may be reduced.
- the hot-rolled steel sheet produced in the hot rolling process may be temper-rolled according to a conventional method, or may be pickled to remove the scale formed on the surface.
- An annealing temperature shall be 700 degreeC and less than 900 degreeC.
- the annealing temperature is 700 ° C. or less, the amount of Ti carbide having a particle diameter of 20 nm or less generated by aging precipitation cannot be sufficiently secured, and the Fe oxide film present on the surface of the hot-rolled steel sheet is not sufficiently removed and good plating is achieved. There is a risk that the sex cannot be secured.
- the annealing temperature is 900 ° C. or higher, the amount of coarse Ti carbide exceeding the particle diameter of 20 nm increases. In any case, desired strength, punching workability and stretch flangeability cannot be obtained.
- the annealing temperature is preferably more than 700 ° C. and less than 880 ° C.
- the holding time for annealing at a temperature exceeding 700 ° C. and less than 900 ° C. is set to be 5 seconds or more and 200 seconds or less.
- the holding time is preferably 10 s or more and 180 s or less.
- the above hot-rolled steel sheet that has undergone the annealing process is plated.
- the plating process may be either electroplating or hot dipping.
- a hot dip galvanizing process may be performed as the plating process, or an alloying process may be further performed after the hot dip galvanizing process.
- the plating bath temperature and the alloying treatment temperature are preferably set to temperatures that do not exceed the annealing temperature from the viewpoint of suppressing excessive coarsening of the Ti carbide.
- the above-mentioned hot rolled galvanized steel sheet may be subjected to temper rolling according to a conventional method.
- Molten steel having the composition shown in Table 1 was melted in a converter, and a slab having a thickness of 250 mm was formed by a continuous casting method. These slabs (steel materials) are heated under the conditions shown in Table 2, and then subjected to the hot rolling process, cooling process, and winding process under the conditions shown in Table 2, and a hot rolled steel sheet having a plate thickness of 2.6 mm and a plate width of 1000 mm. Was made. Subsequently, after pickling and temper rolling with a pressure regulation rate of 0.5%, an annealing process under the conditions shown in Table 2 was performed and immersed in a 450 ° C.
- test pieces were collected from the hot-rolled plated steel sheet obtained as described above, and subjected to a structure observation, a tensile test, a punching test, and a hole expansion test. Moreover, the plating property was evaluated.
- the test method is as follows.
- the number of crystal grains having a grain orientation spread value of 4 ° or less is determined for each crystal grain, the number ratio in the measurement region is determined, and ferrite, tempered bainite Of these, the total area ratio of one or more types was used. At this time, an aggregate of measurement points in which the orientation difference between all adjacent measurement points was less than 5 ° was defined as one crystal grain.
- TEM transmission electron microscope
- Image J image processing software From
- Ti carbide was electrolytically extracted from a test piece collected from the hot-rolled plated steel sheet, the Ti carbide size was fractionated using an alumina filter having a pore diameter of 20 nm, and these were acid-dissolved. Next, the Ti concentration (mass% with respect to all added elements) existing as Ti carbide was determined by inductively coupled plasma emission spectrometry, and the volume fraction of Ti carbide having a particle size of 20 nm or less was determined.
- the area ratio of martensite and retained austenite is determined by first polishing a specimen for observation of structure taken from a hot-rolled steel sheet, polishing a plate thickness section parallel to the rolling direction, and martensite and retentive reagent with a corrosive liquid (repeller reagent). Residual austenite was revealed. Next, the plate thickness 1/4 position and the plate thickness 1/2 position (plate thickness center position) are observed with an optical microscope (magnification: 500 times), and elliptical approximation is performed by image processing software (Image J). Those having a diameter of 1 ⁇ m or more were selected, and the area ratio of martensite and retained austenite having a major axis of 1 ⁇ m or more was determined.
- ⁇ ⁇ ⁇ indicates 5 punch holes with Ra of 15 ⁇ m or less
- ⁇ indicates 3-4 punch holes with Ra of 15 ⁇ m or less
- ⁇ and ⁇ pass. did.
- the punching workability was evaluated with x (failed) other than that.
- Ra exceeds 15 ⁇ m, the chemical conversion property of the punched end face is likely to deteriorate. Therefore, the Ra of the punched end face is 15 ⁇ m or less was used as a pass criterion.
- the plating property of the obtained hot-rolled plated steel sheet was visually evaluated by appearance inspection. The case where the plating layer was formed with the full length and the full width of the hot-rolled plated steel sheet was marked with ⁇ , and the part where the non-plated part was observed was marked with ⁇ .
- a high-strength hot-rolled plated steel sheet having excellent formability with TS of 570 MPa or more, Ra of 15 ⁇ m or less, and ⁇ of 60% or more is obtained.
- the comparative example which is out of the scope of the present invention is inferior in any of TS, Ra, ⁇ , and plating properties.
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Abstract
Description
特許文献2には、質量%で、C:0.02%以上0.075%以下、Si:0.001%以上0.2%以下、Mn:2.0%以上4.5%以下、P:0.1%以下、S:0.01%以下、sol.Al:0.001%以上0.2%以下、N:0.01%以下、O:0.01%以下を含有し、TiおよびNbの添加量が0.14≦Ti+Nb/2≦0.3を満足する成分組成とし、従来の鋼板に比べて多量のMnを添加することにより溶融亜鉛めっき処理工程における粗大なセメンタイト、マルテンサイトやオーステナイト相の生成を抑制し、伸びフランジ性に優れた、引張強さ590MPa以上の高強度溶融亜鉛めっき鋼板を得る方法が開示されている。 特許文献3には、質量%で、C:0.03%以上0.12%以下、Si:0.01%以上0.5%以下、Mn:1.4%以上5.0%以下、P:0.05%以下、S:0.010%以下、sol.Al:0.001%以上0.5%以下およびN:0.020%以下を含有する化学組成の熱延鋼板に、650℃以上950℃以下の温度域まで加熱し、3℃/秒以上20℃/秒以下の平均冷却速度で550℃まで冷却し、420℃以上550℃以下の温度域に20秒間以上90秒間以下保持する熱処理を施すことで、体積率30%以上94%以下のフェライトと、5%以上69%以下のベイナイトならびに残留オーステナイトおよびマルテンサイトを合計で1.0%以上10%以下の金属組織を有する、延性と伸びフランジ性に優れた引張強さ500MPa以上の高強度溶融亜鉛めっき熱延鋼板を得る方法が開示されている。
特許文献4には、質量%で、C:0.02%以上0.10%以下、Si:0.005%以上0.5%以下、Mn:1.4%以上2.5%以下、P:0.025%以下、S:0.010%以下、sol.Al:0.001%以上0.2%以下、N:0.008%以下およびTi:0.15%以下を含有し、さらにCa:0.01%以下、Mg:0.01%以下およびREM:0.01%以下からなる群から選択された1種または2種以上を含有するとともに、下記式(3)~(5)を満足する化学組成を有するとともに、面積%で、フェライト:50~94%、ベイナイト:5~49%ならびにマルテンサイトおよび残留オーステナイトの合計:1~20%を含有する鋼組織を有する、延性と伸びフランジ性に優れた引張強さ:490MPa以上の高強度溶融亜鉛めっき鋼板を得る方法が開示されている。
C-(12/48)×Ti*-(12/93)×Nb≦0.090 ・・・(3)
Ti*=max[Ti-(48/14)×N-(48/32)×S,0] ・・・(4)
2/3×C+(1/150)×Mn+P+2×S<0.15 ・・・(5)
ここで、式(3)~(5)における各元素記号は各元素の含有量(単位:質量%)を示し、式(4)におけるmax[ ]は[ ]内の引数のうち最大の値を返す関数である。
その結果、比較的結晶粒内の転位密度が低いフェライト相や焼戻しベイナイト相を主相とすることで、成形性(打抜き加工性と伸びフランジ性)を向上させることができることを知見した。これは、転位密度が低いことによる母材の靭性向上と、き裂発生の起点となる上記の主相以外の第二相を低減したことによる靭性向上の効果によるものと考えられる。
さらに、Ti炭化物などの金属炭化物の粒子径を20nm以下に制御し、鋼中に微細に分散させることにより、転位密度の低いフェライト相や焼戻しベイナイト相を主相とした場合でも、引張強さ570MPa以上の高強度を安定して得られることが分かった。
加えて、フェライト相や焼戻しベイナイト相を微細な金属炭化物で分散強化することで、主相と第二相の硬度差が小さくなり、伸びフランジ性がさらに向上することも分かった。
本発明は、以上の知見に基づいてなされたものであり、さらに最適な成分組成と金属組織について検討を重ねた結果なされたものである。
[1]質量%で、C:0.03~0.15%、Si:0.4%以下、Mn:1.2~1.9%、Ti:0.05~0.25%、B:0.0005~0.0050%、P:0.03%以下、S:0.005%以下、Al:0.005~0.4%、N:0.01%以下を含有し、残部がFeおよび不可避的不純物からなる成分組成を有し、フェライト、焼戻しベイナイトのうち1種以上が合計で面積率90%以上であり、粒子径が20nm以下であるTi炭化物の体積率が0.05vol%以上である組織を有し、鋼板表面にめっき層又は合金化めっき層を有する高強度熱延めっき鋼板。
[2]前記成分組成に加えて、さらに、質量%で、Nb、V、Zr、Mo、Cr、W、Ta、Hfの1種以上を合計で0.1%以下含有する上記[1]に記載の高強度熱延めっき鋼板。
[3]前記成分組成に加えて、さらに、質量%で、Ca:0.005%以下、Mg:0.005%以下、REM:0.005%以下のうち1種以上を、合計で下記式(1)を満足するように含有する上記[1]または[2]に記載の高強度熱延めっき鋼板。
0.5≦[%X]*/1.25[%S]≦5.0・・・(1)
ここで、[%X]*=[%X]―(0.18+130[%X])×[%O]、[%X]=[%Ca]+[%Mg]+[%REM]、[%O]≦0.005%。
[%S]、[%O]、[%Ca]、[%Mg]、[%REM]は各元素の含有量(質量%)であり、含有しない場合は0とする。
以上のように、本発明は、産業上極めて有用な発明である。
まず、本発明の高強度熱延めっき鋼板の成分組成の限定理由について説明する。なお、以下の成分組成を表す%は、特に断らない限り質量%を意味するものとする。
Cは鋼板中で金属炭化物を形成して強度向上に寄与する元素である。Cの含有量が0.03%未満では所望の強度が得られない。一方、Cの含有量が0.15%を超えると、鋼板中に生成するベイナイト、パーライト、セメンタイト、マルテンサイト、残留オーステナイトの量が増加し、打抜き加工性と伸びフランジ性を悪化させる。このため、C含有量は0.03~0.15%とする。好ましくは0.05~0.13%、より好ましくは0.06~0.13%、さらに好ましくは0.07~0.11%である。
Siは、延性を低下させること無く強度を向上させる元素として有効な元素である。しかし、めっき鋼板においては、めっき層の付着を阻害し鋼板の耐食性を劣化させる。このため、Si含有量の上限を0.4%とする。好ましくは0.2%以下、より好ましくは0.05%以下である。なお、Si含有量はゼロであっても問題ない。
Mnは、固溶強化と結晶粒微細化強化を通じて鋼板の強度向上に寄与する元素である。Mnの含有量が1.2%未満では所望の強度が得られない。一方で、Mnの含有量が1.9%を超えると、焼入れ性が過度に上昇し、粗大なマルテンサイトや残留オーステナイトが生成し、打抜き加工性と伸びフランジ性を悪化させる。このため、Mn含有量は1.2~1.9%とする。好ましくは1.40~1.85%、より好ましくは1.41~1.80%である。
Tiは、めっき処理前に行なう焼鈍処理中に微細なTi炭化物として析出し、鋼板の高強度化に寄与する本発明において重要な元素である。Tiの含有量が0.05%未満では、析出するTi炭化物量が不十分となり、所望の強度が得られない。Tiの含有量が0.25%を超えると、き裂発生の起点となる粗大なTi炭窒化物が増加し、打抜き加工性と伸びフランジ性を悪化させる。このため、Ti含有量は0.05~0.25%とする。好ましくは、0.05~0.20%、より好ましくは0.05~0.15%である。
Bは、焼入れ性を向上させ、ベイナイト、パーライト、セメンタイトの過度な生成を抑制するために添加する。このような効果を得るためには、Bの含有量を0.0005%以上とする必要がある。一方で、Bの含有量が0.0050%を超えると、過度に焼入れ性が向上し、粗大なマルテンサイトや残留オーステナイトが生成し、打抜き加工性と伸びフランジ性を悪化させる。このため、B含有量は0.0005~0.0050%とする。好ましくは、0.0005~0.0040%、より好ましくは0.0005~0.0030%である。
Pは不純物として含有される元素であり、旧オーステナイト粒界に偏析して靭性を低下させるため、打抜き性と伸びフランジ性を高めるためには、極力低減することが好ましいが、0.03%までの含有は許容できる。このため、P含有量は0.03%以下とする。好ましくは0.02%以下である。なお、P含有量はゼロであっても問題ない。
SはTi硫化物を形成し、Ti添加による強度向上効果を阻害する。またMn硫化物を形成し、打抜き性と伸びフランジ性を悪化させる。このため、S含有量を極力低減することが好ましいが、0.005%以下までの含有は許容できる。このため、S含有量は0.005%以下とする。好ましくは、0.004%以下、より好ましくは0.003%以下である。なお、S含有量はゼロであっても問題ない。
Alは、脱酸剤として作用し、鋼の清浄度を向上させるのに有効な元素である。このような効果を得るためにはAlの含有量を0.005%以上とする必要がある。一方、Alの含有量が0.4%を超えると、Al酸化物系介在物の増加を招き、打抜き加工性と伸びフランジ性を悪化させる。このため、Al含有量は0.005~0.4%とする。好ましくは、0.005~0.1%、より好ましくは0.01~0.06%である。
Nは、高温でTiと結合して粗大な窒化物を形成し易く、打抜き加工性と伸びフランジ性を悪化させる。このため、N含有量は0.01%以下とする。好ましくは0.008%以下である。より好ましくは0.006%以下である。なお、N含有量はゼロであっても問題ない。
0.5≦[%X]*/1.25[%S]≦5.0・・・(1)
ここで、[%X]*=[%X]―(0.18+130[%X])×[%O]、
[%X]=[%Ca]+[%Mg]+[%REM]、[%O]≦0.005%
[%S]、[%O]、[%Ca]、[%Mg]、[%REM]は各元素の含有量(質量%)であり、含有しない場合は0とする。
[%X]*/1.25[%S]が0.5を下回ると、Mn硫化物を無害化する効果が無くなる。[%X]*/1.25[%S]が5.0を超えるとCa、Mg、REM酸化物系介在物が増加し、打抜き加工性と伸びフランジ特性を悪化させる場合がある。上記式(1)を満足するように含有すれば、Ca、Mg、REMを添加しても、鋼板中に含有されるOとSとの添加量バランスが適切であるため、粗大な酸化物を形成しこれがき裂の起点となり穴広げ性や打抜き端面性状を悪化させることはない。上記式(1)は、好ましくは0.8以上4.5以下、より好ましくは1.0以上4.0以下である。
本発明の高強度熱延めっき鋼板は、フェライト、焼戻しベイナイトのうち1種以上を主相とし、主相の面積率が合計で90%以上である。また、粒子径が20nm以下であるTi炭化物が微細に分散し粒子径が20nm以下であるTi炭化物の体積率が0.05vol%以上である。
本発明の高強度熱延めっき鋼板は、打抜き加工性および伸びフランジ性を向上させるために、実質的に鋼板の金属組織を転位密度の低い組織のみで構成することが有効である。ゆえに、本発明では、転位密度の低い組織を主相とする。すなわち、本発明において主相とは、フェライトを主相とする場合、焼戻しベイナイトを主相とする場合、フェライトおよび焼戻しベイナイトを主相とする場合のいずれかである。主相(フェライト、焼戻しベイナイトのうち1種以上)の面積率が合計で90%以上であれば、打抜き性や伸びフランジ性を向上させる。好ましくは92%以上である。
本発明において、具体的には、SEM/EBSD法により求まるGrain Orientation Spreadが4°以下となる組織を、転位密度の低い組織と定義する。このGrain Orientation Spreadは一つの結晶粒内のある測定点と同結晶粒内のその他の測定点との方位差を算術平均し、さらに同結晶粒内の全ての測定点について同様の計算を行ない、それらを算術平均することにより求める値である。また、結晶粒内の平均的なひずみ情報、すなわち転位密度と相関する値と考えられる。このようなGrain Orientation Spreadが4°以下となる組織、すなわち、転位密度の低い組織として、熱延工程での相変態により生じるフェライト、ベイナイトを主相とする熱延鋼板をAc3点以下に再加熱した際に生じる焼戻しベイナイト、又は熱延鋼板をAc3点以下に再加熱した際に逆変態オーステナイトから相変態することで生じるフェライト相がある。
本発明では、鋼板に含まれる組織として、上記フェライト、焼戻しベイナイト以外の組織として、ベイナイト、パーライト、マルテンサイト、残留オーステナイト等が挙げられ、これらの組織は合計の面積率で10%以下である。このうち、マルテンサイトと残留オーステナイトに関しては打抜き加工性、伸びフランジ性に及ぼす影響が大きいため、面積率を合計で1%未満、特に長径1μm以上の粗大な形態のものについては、その面積率を合計で1%未満とすることが好ましい。
転位密度の低いフェライトや焼戻しベイナイトを主相とする金属組織において、引張強さ570MPa以上の高強度を得るためには、粒子径20nm以下のTi炭化物を鋼中に分散させ強化を図る必要がある。このようなTi炭化物の微細分散強化にとって、その粒子径は極めて重要であり、粒子径が20nmを超えるTi炭化物が増えると所望の強化量が得られない恐れがある。
ここで、Ti炭化物の粒子径は、透過型電子顕微鏡(TEM)で観察したTi炭化物を、画像解析ソフト(Image J)により円相当径に換算して評価する。
微細分散強化にとって、Ti炭化物の量もまた重要である。粒子径が20nm以下であるTi炭化物の体積率が0.05vol%未満となると、所望の強化量が得られない。ここで、Ti炭化物の体積率fは、式(2)により求められる。
f=(AMC×ρFe×[%MP])/(AM×ρMC)・・・(2)
ここで、AMCは、Ti炭化物の原子量(g/mol)、AMは、Tiの原子量(g/mol)、ρFeは、α鉄の密度(g/cm3)、ρMCは、Ti炭化物の密度(g/cm3)、[%MP]は粒子径20nm以下のTi炭化物中に存在するTiの全添加元素に対する質量濃度を表す。[%MP]は、電解抽出残渣法により20nm以下のTi炭化物を分別し、誘導結合プラズマ発光分析法により求める。
本発明の高強度熱延めっき鋼板は、前記した成分組成からなる鋼素材に、粗圧延と仕上圧延からなる熱間圧延を施し、仕上圧延終了後、冷却し、巻取り、熱延鋼板とし、次いで熱延鋼板を焼鈍し、めっき層を付着させることにより製造することができる。
本発明において、鋼素材の溶製方法は特に限定されず、転炉、電気炉等、公知の溶製方法を採用することができる。また、溶製後、偏析等の問題から、連続鋳造法によりスラブ(鋼素材)とするのが好ましいが、造塊―分塊圧延法、薄スラブ連続鋳造法等の公知の鋳造方法でスラブとしても良い。なお、鋳造後のスラブを熱間圧延するにあたっては、加熱炉でスラブを再加熱した後に圧延しても良いし、スラブが所定温度以上の温度を保持している場合は、再加熱することなく直送圧延しても良い。
仕上圧延後の冷却では、780℃-680℃間の平均冷却速度が30℃/s以上となるように冷却する。平均冷却速度が30℃/s未満であると、冷却中にフェライト変態が進行し、Ti炭化物が析出する。なお、好ましくは60℃/s以上である。
巻取り温度は350℃以上550℃以下とする。巻取り温度が350℃未満では、金属組織が転位密度が極めて高いマルテンサイト相となり、その後の焼鈍工程を経てもGrain Orientation Spreadが4°以下となる転位密度が低い組織を90%以上確保することが困難となる。一方で巻取り温度が550℃を超えると、巻取り工程でフェライト変態が進行し、熱延鋼板中にTi炭化物が析出し、めっき層を付着させる際の焼鈍工程においてTi炭化物が成長し、鋼板強度が低下する恐れがある。
700℃超え900℃未満の焼鈍を施す際の保持時間は5s以上200s以下とする。保持時間が5s未満であると、焼鈍温度を900℃とした場合でも、十分なTi炭化物量を得ることが出来なくなる。一方で、保持時間が200sを超えると、焼鈍温度を700℃とした場合でも、Ti炭化物の粗大化が抑制できない。なお、保持時間は10s以上180s以下とすることが好ましい。
表1に示す成分組成の溶鋼を転炉で溶製し、連続鋳造法で板厚250mmのスラブとした。それらスラブ(鋼素材)を、表2に示す条件で加熱したのち、表2に示す条件の熱延工程、冷却工程、巻取り工程を施し、板厚2.6mm、板幅1000mmの熱延鋼板を作製した。続いて酸洗と調圧率0.5%の調質圧延を行なった後、表2に示す条件の焼鈍工程を施し、450℃の溶融亜鉛浴に浸漬し鋼板表面に亜鉛めっきを形成させた。さらに一部の鋼板では500℃×100秒の条件でめっき層の合金化処理を行なった。
以上により得られた熱延めっき鋼板から試験片を採取し、組織観察、引張試験、打抜き加工試験、穴広げ試験を行なった。また、めっき性を評価した。試験方法は次の通りである。
得られた熱延めっき鋼板の圧延方向に平行な板厚断面を研磨後、コロイダルシリカ溶液を用いて鏡面仕上げ研磨を行ない、走査型電子顕微鏡に備え付けたCCDカメラ(EDAX社製)により、電子線後方散乱回折(Electron backscatter diffraction pattern:EBSD)パターンを取得した。EBSDの測定は、板厚1/4位置および1/2位置からそれぞれ300μm2以上の領域を任意に2視野ずつ選び、電子線の照射間隔(測定間隔)0.5μmの条件で行なった。測定したEBSDパターンをTSL社製の解析ソフトOIM Analysisを用いて、各結晶粒のGrain Orientation Spread値が4°以下である結晶粒の個数を求め、測定領域における個数割合を求め、フェライト、焼戻しベイナイトのうち1種以上の合計の面積率とした。この際、隣接するすべての測定点の方位差が5°未満である測定点の集合体を1つの結晶粒として定義した。
更に、熱延めっき鋼板から採取した薄膜試験片を用いて、日本電子社製の透過型電子顕微鏡(TEM)を用いて、倍率100000倍で100個以上のTi炭化物の観察を行ない、画像処理ソフト(Image J)により、炭化物の円相当粒子径を求めた。
また、熱延めっき鋼板から採取した試験片から、Ti炭化物を電解抽出し、孔径20nmのアルミナフィルタを用いてTi炭化物サイズを分別し、それらを酸溶解した。次に、Ti炭化物として存在するTi濃度(全添加元素に対する質量%)を誘導結合プラズマ発光分析法により求め、粒子径が20nm以下であるTi炭化物の体積率を求めた。
また、マルテンサイトおよび残留オーステナイトの面積率は、先ず、熱延鋼板から採取した組織観察用の試験片を、圧延方向に平行な板厚断面を研磨し、腐食液(レペラ試薬)でマルテンサイトおよび残留オーステナイトを現出させた。次に、板厚1/4位置および板厚1/2位置(板厚中央位置)を光学顕微鏡(倍率:500倍)で観察し、画像処理ソフト(Image J)による楕円近似を行ない、長径が1μm以上のものを選定し、長径が1μm以上のマルテンサイトおよび残留オーステナイトの面積率を求めた。
得られた熱延めっき鋼板から引張方向が圧延方向と直角になるようにJIS 5号引張試験片を採取し、JIS Z 2241の規定に準拠して引張試験を行ない、降伏強さ(YS)、引張強さ(TS)、全伸び(El)を求めた。引張強さ(TS)は570MPa以上、全伸び(EL)は10%以上を合格の基準とした。
得られた熱延めっき鋼板から、30mm角の試験片を採取した。試験片中央部に打抜きポンチを10mmφの平底型として、打抜きクリアランスが10%、15%、20%、25%、30%となるようにダイ側の穴径を選定し、上から板押さえで固定してポンチ穴を打ち抜いた。打ち抜き後、断面が圧延方向と垂直になる打ち抜き端面について、レーザー顕微鏡(株式会社キーエンス製)を用いて板厚方向の算術平均粗さRaを測定した。前記の5個のポンチ穴について、Raが15μm以下であるポンチ穴が5個のものを◎、Raが15μm以下であるポンチ穴が3~4個のものを○とし、◎と○を合格とした。それ以外のものを×(不合格)として、打抜き加工性を評価した。なお、Raが15μmを超えると、打ち抜き端面の化成処理性が悪化しやすくなることから、打ち抜き端面のRaが15μm以下であることを合格の基準として用いた。
得られた熱延めっき鋼板から、100mm角の試験片を採取し、試験片中央部に打抜きポンチを10mmφの平底型として、打抜きクリアランスが12.5%となる条件で打抜き加工し、ポンチ側から頂角60°の円錐ポンチを押し上げて穴を広げた。板厚を貫通する明瞭なき裂が発生した時点で円錐ポンチを止め、その時点の穴直径を測定した。穴広げ後の穴径と穴広げ前の穴径の差を穴広げ前の値で割り、それに100を掛けた数字を穴広げ率(λ)とし、伸びフランジ性の指標とした。穴広げ率(λ)が60%以上を合格とした。
得られた熱延めっき鋼板のめっき性は、外観検査により目視評価した。熱延めっき鋼板の全長、全幅でめっき層が形成されていたものを○、一部不めっき部が観察されたものを×とした。
一方、本発明の範囲を外れる比較例は、TS、Ra、λ、めっき性のうち、いずれかの特性が劣っている。
Claims (3)
- 質量%で、C:0.03~0.15%、
Si:0.4%以下、
Mn:1.2~1.9%、
Ti:0.05~0.25%、
B:0.0005~0.0050%、
P:0.03%以下、
S:0.005%以下、
Al:0.005~0.4%、
N:0.01%以下を含有し、残部がFeおよび不可避的不純物からなる成分組成を有し、
フェライト、焼戻しベイナイトのうち1種以上が合計で面積率90%以上であり、
粒子径が20nm以下であるTi炭化物の体積率が0.05vol%以上である組織を有し、
鋼板表面にめっき層又は合金化めっき層を有する
高強度熱延めっき鋼板。 - 前記成分組成に加えて、さらに、質量%で、Nb、V、Zr、Mo、Cr、W、Ta、Hfの1種以上を合計で0.1%以下含有する請求項1に記載の高強度熱延めっき鋼板。
- 前記成分組成に加えて、さらに、質量%で、Ca:0.005%以下、
Mg:0.005%以下、
REM:0.005%以下のうち1種以上を、合計で下記式(1)を満足するように含有する請求項1または2に記載の高強度熱延めっき鋼板。
0.5≦[%X]*/1.25[%S]≦5.0・・・(1)
ここで、[%X]*=[%X]―(0.18+130[%X])×[%O]、[%X]=[%Ca]+[%Mg]+[%REM]、[%O]≦0.005%。
[%S]、[%O]、[%Ca]、[%Mg]、[%REM]は各元素の含有量(質量%)であり、含有しない場合は0とする。
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WO2010131303A1 (ja) * | 2009-05-11 | 2010-11-18 | 新日本製鐵株式会社 | 打抜き加工性と疲労特性に優れた熱延鋼板、溶融亜鉛めっき鋼板、およびそれらの製造方法 |
JP2011241429A (ja) | 2010-05-17 | 2011-12-01 | Sumitomo Metal Ind Ltd | 溶融亜鉛めっき鋼板およびその製造方法 |
JP2011241456A (ja) | 2010-05-20 | 2011-12-01 | Sumitomo Metal Ind Ltd | 溶融めっき熱延鋼板およびその製造方法 |
JP2012026032A (ja) * | 2010-06-25 | 2012-02-09 | Jfe Steel Corp | 伸びフランジ性に優れた高強度熱延鋼板およびその製造方法 |
JP2013044022A (ja) | 2011-08-24 | 2013-03-04 | Nippon Steel & Sumitomo Metal Corp | 溶融亜鉛めっき鋼板およびその製造方法 |
JP2013100574A (ja) * | 2011-11-08 | 2013-05-23 | Jfe Steel Corp | 材質均一性に優れた高張力熱延鋼板およびその製造方法 |
JP2014148698A (ja) * | 2013-01-31 | 2014-08-21 | Jfe Steel Corp | バーリング加工性に優れた高強度熱延鋼板およびその製造方法 |
JP2014148696A (ja) * | 2013-01-31 | 2014-08-21 | Jfe Steel Corp | バーリング加工性に優れた高強度熱延鋼板およびその製造方法 |
JP2014173151A (ja) | 2013-03-11 | 2014-09-22 | Nippon Steel & Sumitomo Metal | 加工性と疲労特性に優れた高強度熱延鋼板及びその製造方法 |
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CN110402297B (zh) | 2022-04-12 |
US11117348B2 (en) | 2021-09-14 |
KR102286953B1 (ko) | 2021-08-06 |
JPWO2018163871A1 (ja) | 2019-03-14 |
KR20190116359A (ko) | 2019-10-14 |
CN110402297A (zh) | 2019-11-01 |
EP3575429B1 (en) | 2021-08-25 |
EP3575429A1 (en) | 2019-12-04 |
US20210129491A1 (en) | 2021-05-06 |
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