WO2024111525A1 - High-strength hot-rolled steel sheet, and method for producing same - Google Patents
High-strength hot-rolled steel sheet, and method for producing same Download PDFInfo
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- WO2024111525A1 WO2024111525A1 PCT/JP2023/041535 JP2023041535W WO2024111525A1 WO 2024111525 A1 WO2024111525 A1 WO 2024111525A1 JP 2023041535 W JP2023041535 W JP 2023041535W WO 2024111525 A1 WO2024111525 A1 WO 2024111525A1
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- steel sheet
- heating
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- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 102
- 239000010959 steel Substances 0.000 title claims abstract description 102
- 238000004519 manufacturing process Methods 0.000 title claims description 15
- 239000002244 precipitate Substances 0.000 claims abstract description 60
- 229910052758 niobium Inorganic materials 0.000 claims abstract description 49
- 229910052719 titanium Inorganic materials 0.000 claims abstract description 49
- 229910001563 bainite Inorganic materials 0.000 claims abstract description 36
- 229910000734 martensite Inorganic materials 0.000 claims abstract description 33
- 239000002344 surface layer Substances 0.000 claims abstract description 23
- 230000000717 retained effect Effects 0.000 claims abstract description 17
- 239000000203 mixture Substances 0.000 claims abstract description 15
- 229910001566 austenite Inorganic materials 0.000 claims abstract description 12
- 239000012535 impurity Substances 0.000 claims abstract description 6
- 238000001816 cooling Methods 0.000 claims description 51
- 238000010791 quenching Methods 0.000 claims description 12
- 230000000171 quenching effect Effects 0.000 claims description 12
- 229910052718 tin Inorganic materials 0.000 claims description 11
- 238000005098 hot rolling Methods 0.000 claims description 9
- 238000010438 heat treatment Methods 0.000 abstract description 105
- 230000003111 delayed effect Effects 0.000 abstract description 38
- 239000002245 particle Substances 0.000 abstract description 10
- 229910052742 iron Inorganic materials 0.000 abstract description 7
- 239000006104 solid solution Substances 0.000 abstract 2
- 238000012360 testing method Methods 0.000 description 25
- 238000005096 rolling process Methods 0.000 description 21
- 230000007423 decrease Effects 0.000 description 19
- 230000000694 effects Effects 0.000 description 18
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 16
- 229910052761 rare earth metal Inorganic materials 0.000 description 13
- 150000001247 metal acetylides Chemical class 0.000 description 10
- 238000001556 precipitation Methods 0.000 description 10
- 229910000859 α-Fe Inorganic materials 0.000 description 10
- 238000000034 method Methods 0.000 description 9
- 229910052787 antimony Inorganic materials 0.000 description 8
- 239000000463 material Substances 0.000 description 8
- 239000000243 solution Substances 0.000 description 8
- 229910052710 silicon Inorganic materials 0.000 description 7
- XUIMIQQOPSSXEZ-UHFFFAOYSA-N Silicon Chemical compound [Si] XUIMIQQOPSSXEZ-UHFFFAOYSA-N 0.000 description 6
- 239000010703 silicon Substances 0.000 description 6
- 238000005868 electrolysis reaction Methods 0.000 description 5
- 239000003792 electrolyte Substances 0.000 description 5
- 238000009863 impact test Methods 0.000 description 5
- 229910001562 pearlite Inorganic materials 0.000 description 5
- 238000012545 processing Methods 0.000 description 5
- VEXZGXHMUGYJMC-UHFFFAOYSA-N Hydrochloric acid Chemical compound Cl VEXZGXHMUGYJMC-UHFFFAOYSA-N 0.000 description 4
- 229910052802 copper Inorganic materials 0.000 description 4
- 238000005259 measurement Methods 0.000 description 4
- GCLGEJMYGQKIIW-UHFFFAOYSA-H sodium hexametaphosphate Chemical compound [Na]OP1(=O)OP(=O)(O[Na])OP(=O)(O[Na])OP(=O)(O[Na])OP(=O)(O[Na])OP(=O)(O[Na])O1 GCLGEJMYGQKIIW-UHFFFAOYSA-H 0.000 description 4
- 235000019982 sodium hexametaphosphate Nutrition 0.000 description 4
- 238000004458 analytical method Methods 0.000 description 3
- 229910052796 boron Inorganic materials 0.000 description 3
- 239000013078 crystal Substances 0.000 description 3
- 238000004090 dissolution Methods 0.000 description 3
- 229910001568 polygonal ferrite Inorganic materials 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- 238000009864 tensile test Methods 0.000 description 3
- 238000007545 Vickers hardness test Methods 0.000 description 2
- 238000005452 bending Methods 0.000 description 2
- 230000015572 biosynthetic process Effects 0.000 description 2
- 230000000052 comparative effect Effects 0.000 description 2
- 230000007797 corrosion Effects 0.000 description 2
- 238000005260 corrosion Methods 0.000 description 2
- 238000005520 cutting process Methods 0.000 description 2
- 238000000354 decomposition reaction Methods 0.000 description 2
- 238000013461 design Methods 0.000 description 2
- 230000006866 deterioration Effects 0.000 description 2
- 238000001887 electron backscatter diffraction Methods 0.000 description 2
- 239000000446 fuel Substances 0.000 description 2
- 238000007654 immersion Methods 0.000 description 2
- 229910052750 molybdenum Inorganic materials 0.000 description 2
- 229910052759 nickel Inorganic materials 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 238000005498 polishing Methods 0.000 description 2
- 238000003825 pressing Methods 0.000 description 2
- 238000004080 punching Methods 0.000 description 2
- 230000002829 reductive effect Effects 0.000 description 2
- 238000005728 strengthening Methods 0.000 description 2
- 238000005496 tempering Methods 0.000 description 2
- 238000010998 test method Methods 0.000 description 2
- 229910052720 vanadium Inorganic materials 0.000 description 2
- 229910052765 Lutetium Inorganic materials 0.000 description 1
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 1
- 238000002441 X-ray diffraction Methods 0.000 description 1
- 230000001133 acceleration Effects 0.000 description 1
- 239000002253 acid Substances 0.000 description 1
- 239000007864 aqueous solution Substances 0.000 description 1
- 239000003125 aqueous solvent Substances 0.000 description 1
- 229910052785 arsenic Inorganic materials 0.000 description 1
- 229910052799 carbon Inorganic materials 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 229910052804 chromium Inorganic materials 0.000 description 1
- 150000001875 compounds Chemical class 0.000 description 1
- 238000005261 decarburization Methods 0.000 description 1
- 230000003247 decreasing effect Effects 0.000 description 1
- 238000011161 development Methods 0.000 description 1
- 238000005315 distribution function Methods 0.000 description 1
- 239000008151 electrolyte solution Substances 0.000 description 1
- 229910052735 hafnium Inorganic materials 0.000 description 1
- 230000001771 impaired effect Effects 0.000 description 1
- 238000002347 injection Methods 0.000 description 1
- 239000007924 injection Substances 0.000 description 1
- 229910052746 lanthanum Inorganic materials 0.000 description 1
- FZLIPJUXYLNCLC-UHFFFAOYSA-N lanthanum atom Chemical compound [La] FZLIPJUXYLNCLC-UHFFFAOYSA-N 0.000 description 1
- OHSVLFRHMCKCQY-UHFFFAOYSA-N lutetium atom Chemical compound [Lu] OHSVLFRHMCKCQY-UHFFFAOYSA-N 0.000 description 1
- 229910052749 magnesium Inorganic materials 0.000 description 1
- 229910052748 manganese Inorganic materials 0.000 description 1
- 238000004949 mass spectrometry Methods 0.000 description 1
- YLRAQZINGDSCCK-UHFFFAOYSA-M methanol;tetramethylazanium;chloride Chemical compound [Cl-].OC.C[N+](C)(C)C YLRAQZINGDSCCK-UHFFFAOYSA-M 0.000 description 1
- 229910052757 nitrogen Inorganic materials 0.000 description 1
- 230000036961 partial effect Effects 0.000 description 1
- 229910052698 phosphorus Inorganic materials 0.000 description 1
- 239000011148 porous material Substances 0.000 description 1
- 239000000047 product Substances 0.000 description 1
- 230000005855 radiation Effects 0.000 description 1
- 230000000979 retarding effect Effects 0.000 description 1
- 229910052706 scandium Inorganic materials 0.000 description 1
- 238000001878 scanning electron micrograph Methods 0.000 description 1
- 229910052711 selenium Inorganic materials 0.000 description 1
- 238000010008 shearing Methods 0.000 description 1
- 230000003746 surface roughness Effects 0.000 description 1
- 239000001577 tetrasodium phosphonato phosphate Substances 0.000 description 1
- 230000009466 transformation Effects 0.000 description 1
- 230000001131 transforming effect Effects 0.000 description 1
- 229910052721 tungsten Inorganic materials 0.000 description 1
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 1
- 238000004804 winding Methods 0.000 description 1
- 229910052727 yttrium Inorganic materials 0.000 description 1
- 229910052726 zirconium Inorganic materials 0.000 description 1
Classifications
-
- B—PERFORMING OPERATIONS; TRANSPORTING
- B21—MECHANICAL METAL-WORKING WITHOUT ESSENTIALLY REMOVING MATERIAL; PUNCHING METAL
- B21B—ROLLING OF METAL
- B21B1/00—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations
- B21B1/22—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length
- B21B1/24—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process
- B21B1/26—Metal-rolling methods or mills for making semi-finished products of solid or profiled cross-section; Sequence of operations in milling trains; Layout of rolling-mill plant, e.g. grouping of stands; Succession of passes or of sectional pass alternations for rolling plates, strips, bands or sheets of indefinite length in a continuous or semi-continuous process by hot-rolling, e.g. Steckel hot mill
-
- 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
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
- C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C38/00—Ferrous alloys, e.g. steel alloys
-
- C—CHEMISTRY; METALLURGY
- 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
- C22C38/60—Ferrous alloys, e.g. steel alloys containing lead, selenium, tellurium, or antimony, or more than 0.04% by weight of sulfur
Definitions
- the present invention relates to high-strength hot-rolled steel sheets and their manufacturing methods, and in particular to high-strength hot-rolled steel sheets suitable as materials for automotive parts and their manufacturing methods.
- Patent Document 1 discloses a technique for improving stretch flangeability by setting the processing temperature (post-heating temperature) at 400 to 1000°C.
- Patent Document 2 discloses a technique for a high-strength hot-rolled steel sheet with a TS of 730 MPa or more.
- the hot-rolled steel sheet disclosed in Patent Document 2 has a structure in which bainite is the main phase and 80% or more of the total Ti amount is solid-solute Ti. This provides heat treatment hardening with an increase in YS (yield strength) and TS of 100 MPa or more after heat treatment in which the steel sheet is heated to a temperature range of 500°C to the Ac1 transformation point and held for 60 min.
- Patent Document 3 discloses a technique for a hot-rolled steel sheet with excellent delayed fracture resistance and a TS of 120 kgf/mm2 or more .
- Patent Document 1 does not take into account the performance of the part, such as strength and toughness after post-heating, and there is room for improvement.
- the steel structure changes significantly, and when the steel plate (original plate) before post-heating has a high strength of 1180 MPa or more, the impact on the strength becomes significant, so material design that takes into account the impact of post-heating on strength is necessary.
- the steel plate disclosed in Patent Document 2 does not experience a decrease in strength after heat treatment, but on the contrary, it leads to an excessive increase in strength and there is a problem with the toughness after heat treatment due to the significant fine precipitation of carbides, and there is room for improvement.
- Patent Document 3 aims to improve the workability of the original plate and improve the delayed fracture resistance without post-heating.
- the delayed fracture test was evaluated by deep drawing, and no consideration was given to the more severe end surface, or delayed fracture resistance from the end surface processed after post-heating, leaving room for improvement.
- the present invention was made in consideration of the above circumstances, and aims to provide a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
- the inventors focused on the precipitation behavior of Ti and Nb after post-heating of hot-rolled steel sheets, and came up with the idea of improving the properties of the steel sheets after heating (after post-heating) by controlling the initial coarse Ti-containing precipitates and Nb-containing precipitates and the amount of dissolved Ti and Nb before post-heating. Furthermore, they focused on the crystal orientation, and came up with the idea of suppressing delayed fracture after post-heating of the punched end surface by forming a structure in which the surface layer region is concentrated in a specific orientation.
- high strength means that the tensile strength (TS) is 1180 MPa or more and less than 1600 MPa.
- excellent strength after post-heating means that the decrease in strength of the hot-rolled steel sheet after post-heating is 50 or less in Vickers hardness compared to the strength of the hot-rolled steel sheet (original sheet) before post-heating.
- being excellent in toughness after post-heating means that in a Charpy impact test using a test piece taken from the hot-rolled steel sheet after post-heating, the ductile fracture surface ratio at ⁇ 20° C. is 50% or more.
- the thickness of the test piece is 0.6 to 3.0 mm, and when the thickness of the hot-rolled steel sheet exceeds 3.0 mm, the test piece taken from the hot-rolled steel sheet is ground on both sides to a thickness of 3.0 mm, and then subjected to the Charpy impact test.
- excellent resistance to delayed fracture after post-heating means that no cracks are generated when a rectangular test piece having a sheared end surface, which is taken from the steel sheet after post-heating, is V-bent at 90°, the part opened by springback is tightened with a bolt or the like, and the piece is immersed in hydrochloric acid of pH 3 for 96 hours.
- post-heating means a heat treatment in which the hot-rolled steel sheet (original sheet) is heated to 400° C. or higher.
- the present invention has the following configuration.
- C 0.06 to 0.23%, Si: 0.1 to 3.0%, Mn: 1.5 to 3.5%, P: more than 0% and not more than 0.050%; S: more than 0% and 0.0050% or less; Al: more than 0% and not more than 1.5%; N: more than 0% and not more than 0.010%; O: more than 0% and 0.003% or less; Contains Ti and Nb in total 0.040 to 0.200%; The balance is Fe and unavoidable impurities,
- the steel structure has martensite and/or lower bainite as a main phase, and the volume fraction of retained austenite is less than 3%; the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb, that is, (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb), is 0.300 or more and less than 0.800;
- the total amount of Ti and Nb present as precipitates having
- the composition further comprises, in mass%, Cr: 0.005 to 2.0%, Ni: 0.005 to 2.0%, Mo: 0.005 to 1.0%, V: 0.005 to 0.5%, B: 0.0002 to 0.0050%, Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050% Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%
- the high strength hot rolled steel sheet according to [1] comprising one or more selected from the following: [3] A method for producing a high strength hot rolled steel sheet according to the above [1] or [2], A slab having the above-mentioned composition is heated to a temperature range of 1150 to 1300°C and held at that temperature range for 0.2 to 3.5 hours; Next, when hot rolling is performed, The total reduction in the temperature range of 1080°C or more is 80 to 90%, the total reduction in the temperature range of 900°C or less is 20% or more, and the reduction per pass at T
- the steel is allowed to cool for 1.0 s or more.
- the temperature is cooled at an average cooling rate of 50°C/s or more up to 550°C, and the time from reaching 550°C to starting quenching is set to 0.5 to 4.0 s.
- the steel sheet is quenched to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more, and then coiled at the coiling temperature.
- T(°C) 800+1000[Ti]+2500[Nb]
- [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
- the present invention provides a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
- the present invention it is possible to obtain a high-strength hot-rolled steel sheet which is suitable as a material for automobile parts and which has excellent strength, toughness and delayed fracture resistance after post-heating or after post-heat processing.
- a high-strength hot-rolled steel sheet of the present invention it is possible to obtain products such as high-strength automobile parts that exhibit high strength, good toughness, and excellent delayed fracture resistance even after heat treatment is performed to improve workability and fatigue properties.
- the high-strength hot-rolled steel sheet of the present invention may be either a black skin as hot-rolled, or a hot-rolled steel sheet called a white skin which is further pickled after hot rolling.
- the high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 0.6 mm or more.
- the high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 10.0 mm or less. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 1.0 mm or more.
- the thickness is more preferably 6.0 mm or less.
- the width of the high-strength hot-rolled steel sheet of the present invention is preferably 500 mm or more, more preferably 700 mm or more.
- the width of the high-strength hot-rolled steel sheet of the present invention is preferably 1800 mm or less, more preferably 1400 mm or less.
- the high-strength hot-rolled steel sheet of the present invention has a specific chemical composition and a specific steel structure.
- the chemical composition and steel structure will be explained in that order.
- the composition of the high-strength hot-rolled steel sheet of the present invention is, in mass%, C: 0.06-0.23%, Si: 0.1-3.0%, Mn: 1.5-3.5%, P: 0.050% or less (excluding 0%), S: 0.0050% or less (excluding 0%), Al: 1.5% or less (excluding 0%), N: 0.010% or less (excluding 0%), O: 0.003% or less (excluding 0%), Ti and Nb in total 0.040-0.200%, and the balance consisting of Fe and unavoidable impurities.
- C 0.06 to 0.23%
- C is an element that is effective in increasing TS by generating and strengthening martensite and lower bainite, and in suppressing strength reduction after post-heating by combining with Ti, Nb, N, etc. to generate precipitates. If the C content is less than 0.06%, such effects cannot be sufficiently obtained, and the TS of the steel plate (original plate) of 1180 MPa or more or excellent strength after post-heating cannot be obtained. On the other hand, if the C content exceeds 0.23%, the decrease in toughness after post-heating becomes significant, and excellent toughness after post-heating cannot be obtained. Therefore, the C content is set to 0.06 to 0.23%.
- the C content is preferably set to 0.07% or more.
- the C content is preferably set to 0.22% or less, and more preferably set to 0.20% or less.
- Si 0.1 to 3.0% Silicon is an element effective in solution strengthening of steel and suppressing the decrease in strength after post-heating. To obtain such an effect, the silicon content must be 0.1% or more. On the other hand, if the silicon content exceeds 3.0%, polygonal ferrite is excessively formed and the steel structure of the present invention cannot be obtained. Therefore, the silicon content is set to 0.1 to 3.0%.
- the silicon content is preferably set to 0.2% or more.
- the silicon content is preferably set to 2.0% or less, more preferably 1.5% or less.
- Mn 1.5 to 3.5%
- Mn is an element effective in suppressing ferrite and upper bainite and generating lower bainite and martensite. If the Mn content is less than 1.5%, this effect is not sufficiently obtained, and polygonal ferrite, upper bainite, etc. are generated, and the microstructure of the present invention cannot be obtained. On the other hand, if the Mn content exceeds 3.5%, the deterioration of toughness and delayed fracture resistance becomes significant, and excellent toughness and delayed fracture resistance after post-heating cannot be obtained. Therefore, the Mn content is set to 1.5 to 3.5%.
- the Mn content is preferably 1.6% or more.
- the Mn content is preferably 3.0% or less, more preferably 2.5% or less.
- P more than 0% and 0.050% or less P reduces the toughness and delayed fracture resistance after post-heating, so it is desirable to reduce the amount as much as possible.
- a P content of up to 0.050% is acceptable. Therefore, the P content is set to 0.050% or less.
- the P content is preferably set to 0.030% or less. There is no particular lower limit, and the P content may be more than 0%, but if the P content is less than 0.001%, the production efficiency decreases, so the P content is preferably 0.001% or more.
- the S content can be up to 0.0050%. Therefore, the S content is set to 0.0050% or less.
- the S content is preferably set to 0.0030% or less, more preferably set to 0.0020% or less, and even more preferably set to 0.0015%.
- the S content may be more than 0%, but if the S content is less than 0.0002%, the production efficiency decreases, so the S content is preferably 0.0002% or more.
- Al more than 0% and not more than 1.5% Al acts as a deoxidizer, and is preferably added in the deoxidization process.
- the Al content may be more than 0%, but from the viewpoint of using it as a deoxidizer, the Al content is preferably 0.01% or more.
- the Al content is allowed up to 1.5%. Therefore, the Al content is set to 1.5% or less.
- the Al content is preferably set to 0.50% or less, more preferably 0.20% or less.
- N more than 0% and 0.010% or less N generates TiN and NbC and inhibits the precipitation of fine TiC, NbC, etc., so it is preferable to reduce the amount as much as possible.
- an N content of up to 0.010% is acceptable. Therefore, the N content is set to 0.010% or less.
- the N content is preferably set to 0.007% or less. There is no particular lower limit, and the N content may be more than 0%, but if the N content is less than 0.0005%, the production efficiency decreases, so the N content is preferably 0.0005% or more.
- O more than 0% and 0.003% or less O reduces toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible.
- an O content of up to 0.003% is acceptable. Therefore, the O content is set to 0.003% or less.
- the O content is preferably set to 0.002% or less. There is no particular lower limit, and the O content may be more than 0%, but if the O content is less than 0.0002%, production efficiency decreases, so the O content is preferably 0.0002% or more.
- Ti and Nb are the most important elements in the present invention, and are necessary elements for obtaining excellent strength, toughness, and delayed fracture resistance properties after post-heating by generating appropriate fine precipitates such as TiC and NbC after post-heating. If the total content of Ti and Nb is less than 0.040%, such effects are not sufficiently obtained, and excellent strength after post-heating is not obtained. On the other hand, if the total content of Ti and Nb exceeds 0.200%, the amount of coarse precipitates containing Ti and Nb increases, which leads to a decrease in delayed fracture resistance properties after post-heating, and the precipitates after post-heating become excessive, making it impossible to obtain excellent toughness after post-heating.
- the total content of Ti and Nb is set to 0.040 to 0.200%.
- the total content of Ti and Nb is preferably 0.050% or more, more preferably 0.060% or more.
- the total content of Ti and Nb is preferably 0.160% or less, more preferably 0.120% or less.
- the total content of Ti and Nb needs to be within the above range, and the content of either one of them may be 0%.
- the above components are the basic components of the high-strength hot-rolled steel sheet of the present invention.
- the high-strength hot-rolled steel sheet of the present invention contains the above components, with the remainder being Fe and unavoidable impurities.
- the high-strength hot-rolled steel sheet of the present invention can further contain one or more selected from Cr: 0.005-2.0%, Ni: 0.005-2.0%, Mo: 0.005-1.0%, V: 0.005-0.5%, B: 0.0002-0.0050%, Ca: 0.0001-0.0050%, REM: 0.0001-0.0050%, Cu: 0.005-0.5%, Sb: 0.0010-0.10%, Sn: 0.0010-0.10%.
- Cr 0.005 to 2.0% Cr is an element effective in suppressing ferrite and generating lower bainite and martensite.
- the Cr content is preferably 0.005% or more.
- the corrosion resistance may be significantly decreased, so when Cr is contained, the Cr content is preferably 2.0% or less.
- the Cr content is more preferably 0.1% or more.
- the Cr content is more preferably 0.8% or less.
- Ni 0.005 to 2.0%
- Ni is an element effective in suppressing ferrite and generating lower bainite and martensite.
- the Ni content is preferably 0.005% or more.
- the Ni content exceeds 2.0%, a large amount of residual ⁇ is formed, which may lead to a decrease in toughness after post-heating, so when Ni is contained, the Ni content is preferably 2.0% or less.
- the Ni content is more preferably 0.05% or more.
- the Ni content is more preferably 0.8% or less, and further preferably 0.5% or less.
- Mo 0.005 to 1.0%
- Mo is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
- the Mo content is preferably 0.005% or more.
- the Mo content exceeds 1.0%, the generation of Mo-based precipitates becomes significant, which may lead to a decrease in toughness after post-heating, so when Mo is contained, the Mo content is preferably 1.0% or less.
- the Mo content is more preferably 0.05% or more.
- the Mo content is more preferably 0.50% or less.
- V 0.005 to 0.5%
- V is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
- the V content is preferably 0.005% or more.
- the V content exceeds 0.5%, the generation of V-based precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when V is contained, the V content is preferably 0.5% or less.
- the V content is more preferably 0.01% or more.
- the V content is more preferably 0.1% or less.
- B 0.0002 to 0.0050%
- B is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite.
- the B content is preferably 0.0002% or more.
- the B content exceeds 0.0050%, B-based compounds increase, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when B is contained, the B content is preferably 0.0050% or less.
- the B content is more preferably 0.0005% or more.
- the B content is more preferably 0.0040% or less.
- Ca and REM are each an element effective in improving toughness and delayed fracture resistance after post-heating by controlling the shape of inclusions.
- the respective contents are preferably 0.0001% or more.
- the Ca and REM contents are preferably 0.0050% or less.
- the Ca content is more preferably 0.0005% or more.
- the Ca content is more preferably 0.0030% or less.
- the REM content is more preferably 0.0005% or more. Moreover, the REM content is more preferably 0.0030% or less.
- REM is a collective term for Sc, Y, and 15 other elements ranging from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.
- Cu 0.005 to 0.5%
- Sb 0.0010 to 0.10%
- Sn 0.0010 to 0.10%
- Cu, Sb, and Sn are each an element that is effective in retarding the corrosion reaction and improving the delayed fracture resistance after post-heating.
- the Cu content is 0.005% or more
- the Sb content is 0.0010% or more
- the Sn content is 0.0010% or more, respectively.
- the Cu content exceeds 0.5%, the generation of Cu precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when Cu is contained, it is preferable that the Cu content is 0.5% or less.
- the Sb and Sn contents each exceed 0.10%, the grain boundary embrittlement effect becomes excessive, which may lead to a decrease in delayed fracture resistance, so when Sb and Sn are contained, it is preferable that the Sb and Sn contents each are 0.10% or less.
- the Cu content is more preferably 0.05% or more. Furthermore, the Cu content is more preferably 0.3% or less.
- the Sb content is more preferably 0.0050% or more.
- the Sb content is more preferably 0.050% or less.
- the Sn content is more preferably 0.0050% or more.
- the Sn content is more preferably 0.050% or less.
- the effect of the present invention is not impaired even if the content of Cr, Ni, Mo, V, B, Ca, REM, Cu, Sb, and Sn is less than the lower limit value. Therefore, when the content of these components is less than the lower limit value, these elements are treated as unavoidable impurities.
- the present invention may further contain one or more of Mg, As, W, Ta, Pb, Zr, Hf, Te, Bi, and Se in a total amount of 0.3% or less by mass. It is preferable to limit the content of each of these elements to 0.03% or less.
- the steel structure of the high-strength hot-rolled steel sheet of the present invention has martensite and/or lower bainite as the main phase, and the volume fraction of residual ⁇ is less than 3%.
- Main phase martensite and/or lower bainite
- the microstructure in order to obtain high strength and excellent toughness and delayed fracture resistance after post-heating, the microstructure is made to have martensite and/or lower bainite as the main phase. If ferrite, pearlite, residual ⁇ , etc. become the main phase, it becomes difficult to achieve both high strength and excellent toughness and delayed fracture resistance after post-heating. Therefore, the steel microstructure is made to have martensite and/or lower bainite as the main phase.
- the martensite may be either auto-tempered martensite or tempered martensite, but fresh martensite having no carbides inside is excluded.
- the lower bainite may be tempered lower bainite.
- the main phase means a phase that occupies 50% or more in terms of area ratio.
- the area ratio of the main phase is preferably 60% or more, and more preferably 75% or more.
- martensite may be the main phase
- lower bainite may be the main phase
- the total of martensite and lower bainite may be the main phase.
- the upper limit of the area ratio of the main phase is not particularly limited and may be 100%.
- the area ratio of the main phase may be, for example, less than 100% or 98% or less.
- Amount of retained austenite (residual ⁇ ) less than 3% Since retained austenite (residual ⁇ ) is a structure that significantly reduces strength and toughness by transforming into pearlite after post-heating, it is preferable to reduce it as much as possible.
- the volume fraction of retained ⁇ is allowed to be less than 3%. Therefore, the volume fraction of retained ⁇ is set to less than 3%.
- the volume fraction of retained ⁇ is preferably less than 2%, and more preferably less than 1%. There is no particular limit on the lower limit of the volume fraction of retained ⁇ , and the volume fraction of retained ⁇ may be 0%.
- the phases other than martensite, lower bainite, and residual gamma may be one or more of ferrite, pearlite, and upper bainite.
- the total area ratio of the other phases is preferably 30% or less, and more preferably 25% or less. There is no particular lower limit to the area ratio of the other phases, and the total area ratio of the other phases may be 0%.
- (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) is 0.300 or more and less than 0.800. It is preferably 0.350 or more. It is also preferably 0.700 or less.
- the value of (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb) is determined by the method described in the Examples.
- the growth of the precipitates competes with the precipitation of new TiC, NbC, etc. during post-heating. This appropriately suppresses the precipitation of fine TiC, NbC, etc., and can suppress excessive strength increase and toughness decrease.
- the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 mass% or more.
- the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.030 mass% or less. Therefore, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 to 0.030 mass%. Preferably, it is 0.013 mass% or more. Also, it is preferably 0.027 mass% or less.
- the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is determined by the method described in the Examples.
- Pole density of ⁇ 110 ⁇ 111> orientation in the surface layer region from the surface to 100 ⁇ m in the center direction of the sheet thickness 1.8 to 5.0
- the surface layer region from the surface of the steel plate to 100 ⁇ m in the thickness center direction strongly influences the formation of fracture surface during punching or high-speed deformation.
- the pole density of the ⁇ 110 ⁇ 111> orientation in this region to the range of 1.8 to 5.0, excellent toughness can be obtained after post-heating, and the fracture surface properties of punching become good, and excellent delayed fracture resistance properties can be obtained after post-heating.
- the pole density of the ⁇ 110 ⁇ 111> orientation needs to be 1.8 or more in the surface layer region from the surface to 100 ⁇ m in the thickness center direction.
- the pole density of the ⁇ 110 ⁇ 111> orientation is set to 1.8 to 5.0 in the surface layer region from the surface to 100 ⁇ m in the thickness center direction. It is preferably set to 2.0 or more. It is also preferably set to 4.0 or less, and more preferably set to 3.0 or less.
- the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region extending from the surface to 100 ⁇ m in the sheet thickness center direction is determined by the method described in the examples.
- the high-strength hot-rolled steel sheet of the present invention is produced by heating a slab having the above-mentioned composition to a temperature range of 1150 to 1300°C, holding the slab in the temperature range for 0.2 to 3.5 hours, and then hot rolling the slab under conditions in which the total reduction in the temperature range of 1080°C or higher is 80 to 90%, the total reduction in the temperature range of 900°C or lower is 20% or more, and the reduction per pass at or below T (°C) calculated by the following formula is 25% or less, followed by allowing the slab to cool for 1.0 s or more, then cooling the slab at a temperature range up to 550°C at an average cooling rate of 50°C/s or more, setting the time from reaching 550°C to the start of quenching to 0.5 to 4.0 s, then quenching the slab to a coiling temperature of 100 to 250°C at a cooling rate of 200°C/s or more, and
- T(°C) 800+1000[Ti]+2500[Nb]
- [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
- the total reduction in the temperature range of 1080° C. or higher is determined from the ratio of the slab thickness before hot rolling to the plate thickness at 1080° C.
- the total reduction in the temperature range of 900° C. or lower is determined from the ratio of the plate thickness at 900° C. to the final plate thickness.
- the reduction per pass at or below T (° C.) is determined from the ratio of the plate thickness before and after each pass of rolling at or below T (° C.).
- the above temperatures are the surface temperatures at the center of the width of the steel plate, and the above average cooling rates and cooling speeds are the average cooling rate and cooling speed at the surface at the center of the width of the steel plate, respectively. Furthermore, unless otherwise specified, the average cooling rate is [(cooling start temperature - cooling stop temperature) / cooling time from cooling start temperature to cooling stop temperature].
- the heating temperature of the slab is set to 1150 to 1300°C.
- the heating temperature is preferably 1170° C. or higher, and more preferably 1185° C. or higher.
- the heating temperature is preferably 1280° C. or lower, and more preferably 1265° C. or lower.
- Holding time in the temperature range of 1150 to 1300°C 0.2 to 3.5 hours If the holding time in the temperature range of 1150 to 1300°C is less than 0.2 hours, the dissolution of Ti-containing precipitates and Nb-containing precipitates will be insufficient. As a result, a value of (amount of dissolved Ti + amount of dissolved Nb)/(total amount of Ti + total amount of Nb) of 0.300 or more and less than 0.800, or a value of 0.010 to 0.030 mass% of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more will not be obtained.
- the holding time of the slab in the above temperature range is set to 0.2 to 3.5 hours.
- the holding time is preferably 0.4 hours or more.
- the holding time is preferably 2.5 hours or less.
- Total reduction rate at temperatures above 1080°C: 80-90% By carrying out a total reduction of 80 to 90% in a temperature range of 1080°C or more, it is possible to promote the generation and growth of coarse Ti-containing precipitates and Nb-containing precipitates having a particle size of 100 nm or more. As a result, the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more can be set to 0.010 to 0.030 mass%. If the total reduction is less than 80%, the generation of precipitates having a particle size of 100 nm or more is insufficient, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more is less than 0.010 mass%.
- the total reduction in a temperature range of 1080°C or more is set to 80 to 90%.
- the total reduction is preferably set to 81% or more.
- the total rolling reduction is preferably 88% or less.
- Total rolling reduction in the temperature range of 900°C or less is 20% or more If the total rolling reduction in the temperature range of 900°C or less is less than 20%, strain-induced precipitation is suppressed, Ti-containing precipitates and Nb-containing precipitates are reduced, and a value of (solubilized Ti amount + solid-solubilized Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. Or, the texture of the surface layer portion is insufficiently developed, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region is not able to be 1.8 to 5.0. Therefore, the total rolling reduction in the temperature range of 900°C or less is set to 20% or more.
- the upper limit of the total rolling reduction is not particularly limited, but the total rolling reduction is preferably 80% or less, more preferably 60% or less.
- T (°C) or less When reduction of more than 25% per pass is applied at T (°C) or less calculated by the following formula, strain-induced precipitation is promoted, Ti-containing precipitates and Nb-containing precipitates increase, and a value of (solute Ti amount + solute Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained.
- the texture of the surface layer part develops, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. Therefore, the reduction rate per pass at T (°C) or less is set to 25% or less.
- the reduction rate is preferably set to 20% or less, more preferably set to 18% or less.
- the lower limit of the reduction rate is not particularly limited, but since coarse grains may occur at 5% or less, the reduction rate is preferably set to more than 5%.
- the reduction rate is more preferably set to 7% or more.
- [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
- Cooling for 1.0 s or more By cooling after rolling under the above conditions, partial strain is released, strain-induced precipitation and dislocation precipitation during subsequent cooling are suppressed, and Ti-containing precipitates and Nb-containing precipitates can be reduced. To obtain such an effect, it is necessary to set the cooling time after rolling to 1.0 s or more.
- the cooling time is preferably 1.5 s or more, more preferably 2.0 s or more, and even more preferably 2.2 s or more.
- There is no particular limit to the upper limit of the cooling time but if the cooling time is 5.0 s or less, it becomes easier to control the subsequent hot rolling, so the cooling time is preferably 5.0 s or less.
- cooling means exposure to the atmosphere (air cooling) without active cooling (accelerated cooling) by water injection or the like.
- hot rolling includes rough rolling and finish rolling, and the cooling time after rolling is the cooling time after hot rolling, i.e., after finish rolling.
- Cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C After the above cooling, cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C. If the average cooling rate up to 550°C is less than 50°C/s, excessive generation of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc., and formation of the crystal orientation in the surface layer region will be caused. As a result, the phase structure of the present invention, the precipitates, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 will not be obtained.
- the average cooling rate in the temperature range from the cooling start temperature to 550°C after the above cooling is set to 50°C/s or more.
- the average cooling rate is preferably set to 70°C/s or more.
- the average cooling rate is preferably less than 500°C/s, and more preferably less than 200°C/s.
- Time from reaching 550°C to starting quenching 0.5 to 4.0 s
- a certain time leaving a certain time interval
- starting quenching quenching at a cooling rate of 200°C/s or more, which will be described later
- bainite can be formed in the medium temperature region near the surface layer.
- the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of the present invention can be obtained. If the time from reaching 550°C to starting quenching is less than 0.5 s, such an effect cannot be obtained sufficiently, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained.
- the time from 550°C to starting quenching is set to 0.5 to 4.0 s.
- the time is preferably 0.7 s or more.
- the time is preferably 2.0 s or less, more preferably 1.6 s or less.
- Cooling rate to coiling temperature of 100 to 250°C: 200°C/s or more quenching is started after a time of 0.5 to 4.0 s has elapsed between reaching 550°C and starting quenching. If the cooling (quenching) rate to coiling temperature of 100 to 250°C is less than 200°C/s, upper bainite and residual ⁇ are excessively generated, and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region is increased. As a result, the phase structure of the present invention and the pole density of the ⁇ 110 ⁇ 111> orientation in the surface layer region cannot be obtained. Therefore, the cooling rate to the coiling temperature is set to 200°C/s or more. The cooling rate is preferably set to 250°C/s or more. The upper limit of the cooling rate is not particularly limited, but from the viewpoint of shape stability, the cooling rate is preferably 1000°C/s or less, and more preferably 500°C/s or less.
- Winding temperature 100 to 250°C
- the coiling temperature is set to 100 to 250°C.
- the coiling temperature is preferably 120°C or higher.
- the coiling temperature is preferably 220°C or lower.
- the high-strength hot-rolled steel sheet of the present invention has excellent strength, toughness, and delayed fracture resistance after post-heating.
- the heating temperature of the post-heating is 400°C or higher.
- the upper limit of the heating temperature of the post-heating is not particularly limited, but an example of the heating temperature of the post-heating is 1150°C or lower.
- the heating time of the post-heating (holding time at the heating temperature) is not particularly limited, but an example of the heating time is more than 0 seconds.
- the heating time of the post-heating is 3600 seconds or less, for example.
- the hot-rolled steel sheet was used to observe the structure, analyze solute Ti, solute Nb, Ti-containing precipitates, and Nb-containing precipitates, and evaluate the tensile properties according to the following test methods. Furthermore, the hot-rolled steel sheet was post-heated as shown in Table 2, and the hardness, toughness, and delayed fracture resistance were evaluated according to the following test methods using the hot-rolled steel sheet after post-heating.
- the post-heating temperature was 400°C or higher, at which an improvement in stretch flangeability is observed, and the post-heating time was 3600 s or less from the viewpoint of productivity.
- the area ratio of martensite and lower bainite refers to the ratio of the area of each structure to the observed area.
- the area ratio of martensite was measured by cutting a sample from the obtained hot-rolled steel sheet, polishing the plate thickness cross section parallel to the rolling direction, corroding it with 3% nital, and photographing the plate thickness 1/4 position at a magnification of 1500 times with a SEM (scanning electron microscope) in three fields of view.
- the area ratio of each structure was calculated from the image data of the obtained secondary electron image using Image-Pro manufactured by Media Cybernetics, and the average area ratio of the three fields of view was taken as the area ratio of each structure.
- the structure may be determined by a general classification, but can be determined, for example, as follows.
- lower bainite is distinguished as black or dark gray, gray, or light gray containing oriented carbides.
- Martensite is a structure of black to light gray that is regular but contains carbides of multiple orientations. Alternatively, it is observed as white or light gray without carbides.
- the retained austenite is observed as white or light gray without containing carbides. Since it may be difficult to distinguish between a part of martensite and the retained austenite, the retained austenite was obtained by the method described below, and the area ratio of the martensite was obtained by subtracting it from the total area ratio of the martensite and the retained austenite obtained from the SEM image.
- the carbides are white dots or lines.
- ferrite is a structure that is black or dark gray and does not have a substructure such as carbides or laths inside
- pearlite can be distinguished as a black and white layered or partially interrupted layered structure.
- upper bainite can be distinguished as a structure that is black or dark gray and has a substructure such as carbides or laths inside.
- the amount of retained ⁇ is obtained as follows.
- the hot-rolled steel sheet was ground to 1/4+0.1 mm of the sheet thickness, and then chemically polished to a further 0.1 mm to obtain the measurement surface.
- the measurement surface was measured using an X-ray diffraction apparatus with Mo K ⁇ 1 radiation to measure the integral reflection intensities of the (200), (220), and (311) surfaces of fcc iron (austenite) and the (200), (211), and (220) surfaces of bcc iron (ferrite).
- the volume fraction was then calculated from the intensity ratio of the integral reflection intensity from each surface of the fcc iron to the integral reflection intensity from each surface of the bcc iron, and this was taken as the amount of residual ⁇ .
- Table 3 shows the structures constituting the main phase and other structures that account for 50% or more of the area ratio of each structure obtained.
- M means martensite
- LB means lower bainite
- ⁇ means retained austenite
- O means other phases.
- the other phases include one or more of ferrite, pearlite, and upper bainite.
- Pole density of ⁇ 110 ⁇ 111> orientation in the surface region from the surface to 100 ⁇ m in the thickness center direction Samples were cut out from the obtained hot-rolled steel sheet, the thickness cross section parallel to the rolling direction was polished, and strain was removed by electrolytic polishing. After that, crystal orientation data was obtained for the surface region from the surface to 100 ⁇ m in the thickness center direction by EBSD (electron backscatter diffraction). Each sample was measured in three fields of view with a measurement area of 100 ⁇ m x 100 ⁇ m, an acceleration voltage of 30 kV, and a step size of 100 nm. OIM Analysis Ver. 7.3.0 manufactured by TSL Solutions was used to analyze the obtained data.
- the average value of the pole densities of the three fields of view of each sample was taken as the pole density of each sample.
- the concentration ratios of Ti and Nb to Fe were calculated, and further multiplied by the content (mass%) of Fe in the test piece to obtain the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%).
- the content (mass%) of Fe in the test piece was obtained by subtracting the total content (mass%) of components other than Fe from 100 mass%.
- the ratio of the total dissolved Ti amount (mass%) and dissolved Nb amount (mass%) to the total contained Ti amount (mass%) and contained Nb amount (mass%) was calculated.
- the test piece having the precipitate attached to the surface after electrolysis was taken out from the electrolytic solution and immersed in an aqueous solution of sodium hexametaphosphate (500 mg/L) (hereinafter referred to as an aqueous SHMP solution). Then, ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution. Next, the aqueous SHMP solution containing the precipitate was filtered using a filter with a pore size of 100 nm, and the precipitate collected on the 100 nm filter was decomposed with acid, and the decomposition solution was analyzed using an ICP emission spectrometer, and the absolute values of Ti and Nb in the decomposition solution were measured.
- an aqueous SHMP solution sodium hexametaphosphate (500 mg/L)
- ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution.
- the absolute values of Ti and Nb obtained were divided by the amount of electrolyte to obtain the amount of Ti and the amount of Nb (mass %) contained in the precipitates having a particle size of 100 nm or more when the total composition of the test piece was taken as 100 mass %).
- the total of the obtained amount of Ti (mass %) and amount of Nb (mass %) was divided by the total amount of Ti (mass %) and amount of Nb (mass %) contained in the test piece to obtain the total amount of Ti (mass %) present as precipitates containing Ti having a particle size of 100 nm or more and the amount of Nb (mass %) present as precipitates containing Nb having a particle size of 100 nm or more.
- the amount of electrolyte was obtained by measuring the mass of the test piece after the precipitates were peeled off and subtracting it from the mass of the test piece before electrolysis.
- JIS No. 5 tensile test pieces (JIS Z 2241:2011) were taken from the obtained hot-rolled steel sheets in the direction parallel to the rolling direction, and a tensile test was carried out in accordance with the provisions of JIS Z 2241:2011 at a strain rate of 10 ⁇ 3 /s to determine TS. In the present invention, a TS of 1180 MPa or more was considered to be acceptable.
- Vickers hardness test Samples were cut out from the obtained hot-rolled steel sheet and the hot-rolled steel sheet after post-heating, and the cross section of the sheet thickness parallel to the rolling direction was polished. Then, a Vickers hardness test was performed at 1/4 of the sheet thickness position with a load of 5 kg and five measurement points, and the average (arithmetic mean) was taken as the Vickers hardness of the steel sheet. A difference in hardness ( ⁇ HV) of 50 or less before and after post-heating was judged to be excellent in strength after post-heating and was considered to have passed the test.
- ⁇ HV hardness
- Charpy impact test From the hot-rolled steel sheet obtained by post-heating the hot-rolled steel sheet, a test piece with a width of 10 mm and a length of 55 mm was taken, and a V-notch with a tip angle of 45°, a tip radius of 0.25 mm, and a depth of 2 mm was made to prepare a Charpy impact test piece. Then, in accordance with JIS Z 2242:2018, a Charpy impact test was performed five times at -20 ° C. to evaluate the ductile fracture rate. A test piece with an average ductile fracture rate of 50% or more after five tests was judged to have excellent toughness after post-heating and was passed. The plate thickness was 2.9 mm, and the notch direction was parallel to the rolling direction.
- Delayed fracture test From the obtained hot-rolled steel sheet, a test piece with a width of 30 mm and a length of 110 mm was taken, and the post-heating treatment shown in Table 2 was performed to obtain a test piece. This was subjected to 90° V-bending with a bending radius of 15 mm so that the ridge line was parallel to the rolling direction, and bolts were tightened by the amount of opening due to springback, and the test piece was immersed in hydrochloric acid of pH 3 for 96 hours to check for the presence or absence of cracks. Those that did not have cracks were judged to have excellent delayed fracture resistance after post-heating and were passed.
- the end faces of the test pieces were formed by shearing with a shear angle of 1° and a clearance of 10%, and burrs were bent on the outside.
- the "delayed fracture time (hr)” in Table 3 indicates the time when cracks occurred in the test piece. However, “96” in the “delayed fracture time (hr)” indicates that no cracks occurred in the test piece after the above 96-hr immersion.
- All of the inventive examples have a TS of 1180 MPa or more, and are excellent in strength, toughness, and delayed fracture resistance after post-heating.
- the comparative examples that fall outside the scope of the present invention either do not have the desired strength (TS) or do not achieve one or more of the desired strength, toughness, and delayed fracture resistance after post-heating.
- the present invention it is possible to obtain a high-strength hot-rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and excellent strength, toughness and delayed fracture resistance after post-heating.
- the high-strength steel sheet of the present invention is used for automobile parts, it can greatly contribute to improving the collision safety and fuel efficiency of automobiles.
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Abstract
Provided is a high-strength hot-rolled steel sheet having excellent post-heating strength, toughness, and delayed fracture resistance. The high-strength hot-rolled steel sheet has a component composition containing, by mass%, 0.06-0.23% of C, 0.1-3.0% of Si, 1.5-3.5% of Mn, 0-0.050% (exclusive of 0) of P, 0-0.0050% (exclusive of 0) of S, 0-1.5% (exclusive of 0) of Al, 0-0.010% (exclusive of 0) of N, 0-0.003% (exclusive of 0) of O, and 0.040-0.200% of a total of Ti and Nb, with the balance consisting of Fe and unavoidable impurities, wherein: the steel structure has martensite and/or lower bainite as a main phase; the retained austenite is less than 3% by volume; (Ti in solid solution + Nb in solid solution)/(total Ti amount + total Nb amount) is 0.300-0.800 (exclusive of 0.800); the total amount of Ti and Nb present as precipitates having a particle diameter of at least 100 nm is 0.010-0.030 mass%; and the [110]<111> orientation pole density in the surface layer region is 1.8-5.0.
Description
本発明は、高強度熱延鋼板及びその製造方法に関し、特に、自動車用部品の素材として好適な、高強度熱延鋼板及びその製造方法に関するものである。
The present invention relates to high-strength hot-rolled steel sheets and their manufacturing methods, and in particular to high-strength hot-rolled steel sheets suitable as materials for automotive parts and their manufacturing methods.
自動車の衝突安全性改善と燃費向上の観点から、自動車用部品に用いられる熱延鋼板には、高強度化が求められている。一方で、高強度化した熱延鋼板では、プレス時に加工性不足に起因した割れ発生が顕著となるため、プレス工法や鋼板の加工性の改善が必要とされる。工法に関しては、鋼板（原板）を加熱することで加工性の改善を図る検討がされている。なお、本明細書において、鋼板（原板）を部品に加工等する際に加える加熱処理を、後加熱ともいう。また一方で、鋼板に関しては、後加熱して加工した後（後加熱加工後）の部品の特性を考慮した開発が行われている。さらに引張強さ（ＴＳ）が１１８０ＭＰａ以上の高強度鋼板では遅れ破壊も問題となってくるため、後加熱後の耐遅れ破壊特性も考慮した素材設計が必要となる。
In order to improve the crashworthiness and fuel economy of automobiles, there is a demand for higher strength in hot-rolled steel sheets used in automobile parts. On the other hand, in high-strength hot-rolled steel sheets, cracks caused by insufficient workability become prominent during pressing, so improvements in pressing methods and workability of steel sheets are required. Regarding the method, studies are being conducted to improve workability by heating steel sheets (original sheets). In this specification, the heat treatment applied when processing steel sheets (original sheets) into parts is also called post-heating. On the other hand, regarding steel sheets, development is being conducted taking into account the characteristics of parts after post-heating and processing (after post-heating processing). Furthermore, delayed fracture is also an issue for high-strength steel sheets with a tensile strength (TS) of 1180 MPa or more, so material design that takes into account delayed fracture resistance after post-heating is also required.
特許文献１には、加工温度（後加熱の加熱温度）を４００～１０００℃とすることで伸びフランジ性を改善した工法に関する技術が開示されている。特許文献２には、ＴＳが７３０ＭＰａ以上の高強度熱延鋼板に関する技術が開示されている。特許文献２に開示された熱延鋼板では、ベイナイトを主相とし、全Ｔｉ量の８０％以上を固溶Ｔｉとした組織とする。これにより、５００℃～Ａｃ１変態点の温度域に加熱し６０ｍｉｎ間保持する熱処理を施した後のＹＳ（降伏強さ）およびＴＳ増加量が１００ＭＰａ以上である熱処理硬化性が得られる。特許文献３には、１２０ｋｇｆ／ｍｍ２以上のＴＳを有する、耐遅れ破壊特性に優れた熱延鋼板に関する技術が開示されている。
Patent Document 1 discloses a technique for improving stretch flangeability by setting the processing temperature (post-heating temperature) at 400 to 1000°C. Patent Document 2 discloses a technique for a high-strength hot-rolled steel sheet with a TS of 730 MPa or more. The hot-rolled steel sheet disclosed in Patent Document 2 has a structure in which bainite is the main phase and 80% or more of the total Ti amount is solid-solute Ti. This provides heat treatment hardening with an increase in YS (yield strength) and TS of 100 MPa or more after heat treatment in which the steel sheet is heated to a temperature range of 500°C to the Ac1 transformation point and held for 60 min. Patent Document 3 discloses a technique for a hot-rolled steel sheet with excellent delayed fracture resistance and a TS of 120 kgf/mm2 or more .
しかしながら、特許文献１では、後加熱後の強度や靭性等の部品としてのパフォーマンスを考慮しておらず、改善の余地がある。特に４００℃を超えるような高温で後加熱を施すと鋼組織が大きく変化し、後加熱前の鋼板（原板）が１１８０ＭＰａ以上の高強度となるとその強度への影響は顕著となるため、後加熱による強度への影響を考慮した素材設計が必要となる。特許文献２に開示された鋼板は、熱処理後の強度低下は生じないものの、逆に過度な強度上昇を招くとともに炭化物の微細析出が顕著なため熱処理後の靭性に課題があり、改善の余地がある。また、特許文献２では、検討されている鋼板の強度レベルも９８０ＭＰａ程度に留まり、１１８０ＭＰａ以上の高強度における知見や示唆は無い。特許文献３は、原板での加工性を改善し、後加熱無しでの耐遅れ破壊特性向上を志向したものである。しかしながら、特許文献３では、遅れ破壊試験は深絞りで評価されており、より厳しい端面、ひいては後加熱後に加工した端面からの耐遅れ破壊特性についてはなんら考慮されておらず、改善の余地がある。
However, Patent Document 1 does not take into account the performance of the part, such as strength and toughness after post-heating, and there is room for improvement. In particular, when post-heating is performed at a high temperature exceeding 400°C, the steel structure changes significantly, and when the steel plate (original plate) before post-heating has a high strength of 1180 MPa or more, the impact on the strength becomes significant, so material design that takes into account the impact of post-heating on strength is necessary. The steel plate disclosed in Patent Document 2 does not experience a decrease in strength after heat treatment, but on the contrary, it leads to an excessive increase in strength and there is a problem with the toughness after heat treatment due to the significant fine precipitation of carbides, and there is room for improvement. In addition, the strength level of the steel plate examined in Patent Document 2 is only about 980 MPa, and there is no knowledge or suggestion for high strength of 1180 MPa or more. Patent Document 3 aims to improve the workability of the original plate and improve the delayed fracture resistance without post-heating. However, in Patent Document 3, the delayed fracture test was evaluated by deep drawing, and no consideration was given to the more severe end surface, or delayed fracture resistance from the end surface processed after post-heating, leaving room for improvement.
本発明は、上記事情に鑑みてなされたものであり、後加熱後の強度、靭性および耐遅れ破壊特性に優れる高強度熱延鋼板及びその製造方法を提供することを目的とする。
The present invention was made in consideration of the above circumstances, and aims to provide a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
本発明者らは、熱延鋼板の後加熱後のＴｉおよびＮｂの析出挙動に着目し、後加熱前の初期の粗大なＴｉ含有析出物およびＮｂ含有析出物と固溶Ｔｉ量および固溶Ｎｂ量を制御することによる加熱後（後加熱後）の鋼板特性の向上を想到した。さらに、結晶方位にも着目し、表層領域を特定の方位に集積させた組織とすることで、打抜き端面の後加熱後の遅れ破壊を抑制することを想到した。その結果、化学成分を調整したうえでマルテンサイトおよび／または下部ベイナイトを主相とし、残留オーステナイト（残留γ）量を体積率で３％未満とし、表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度を１．８～５．０とし、Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）を０．３００以上０．８００未満、かつ、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量を合計で０．０１０～０．０３０質量％とすることで熱延鋼板（原板）で高強度を発現しつつ、後加熱後も原板に近い強度と、優れた靭性および耐遅れ破壊特性が得られることを見出し、本発明を完成するに至った。
The inventors focused on the precipitation behavior of Ti and Nb after post-heating of hot-rolled steel sheets, and came up with the idea of improving the properties of the steel sheets after heating (after post-heating) by controlling the initial coarse Ti-containing precipitates and Nb-containing precipitates and the amount of dissolved Ti and Nb before post-heating. Furthermore, they focused on the crystal orientation, and came up with the idea of suppressing delayed fracture after post-heating of the punched end surface by forming a structure in which the surface layer region is concentrated in a specific orientation. As a result, by adjusting the chemical composition, it was found that by making martensite and/or lower bainite the main phase, making the amount of retained austenite (residual gamma) less than 3% by volume, making the pole density of the {110}<111> orientation 1.8 to 5.0 in the surface layer region from the surface to 100 μm in the sheet thickness center direction, making the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb (Solute Ti + Solute Nb)/(Total Ti + Total Nb) 0.300 or more and less than 0.800, and making the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more 0.010 to 0.030 mass%, it is possible to obtain high strength in the hot-rolled steel sheet (original sheet), strength close to that of the original sheet even after post-heating, and excellent toughness and delayed fracture resistance, and thus completed the present invention.
なお、本発明において、高強度とは、引張強さ（ＴＳ）が１１８０ＭＰａ以上１６００ＭＰａ未満であることを意味する。
また、本発明において、後加熱後の強度に優れるとは、後加熱後の熱延鋼板の強度の低下が、後加熱前の熱延鋼板（原板）の強度に対して、ビッカース硬さで５０以下であることを意味する。
本発明において、後加熱後の靭性に優れるとは、後加熱後の熱延鋼板から採取した試験片を用いたシャルピー衝撃試験において、－２０℃での延性破面率が５０％以上であることを意味する。なお、前記試験片の板厚は０．６～３．０ｍｍとし、熱延鋼板の板厚が３．０ｍｍを超える場合は、熱延鋼板から採取した試験片を板厚３．０ｍｍまで表裏面研削加工し、シャルピー衝撃試験に供するものとする。
本発明において、後加熱後の耐遅れ破壊特性に優れるとは、後加熱後の鋼板から採取した、せん断端面を有する短冊試験片に９０゜でＶ曲げ加工を施し、次いでスプリングバックで開いた分をボルトなどで締め付け、ｐＨ３の塩酸に９６ｈｒ浸漬した時に割れが発生しないことを意味する。
なお、本発明において、後加熱とは、熱延鋼板（原板）を４００℃以上に加熱する熱処理を意味する。 In the present invention, high strength means that the tensile strength (TS) is 1180 MPa or more and less than 1600 MPa.
In the present invention, excellent strength after post-heating means that the decrease in strength of the hot-rolled steel sheet after post-heating is 50 or less in Vickers hardness compared to the strength of the hot-rolled steel sheet (original sheet) before post-heating.
In the present invention, being excellent in toughness after post-heating means that in a Charpy impact test using a test piece taken from the hot-rolled steel sheet after post-heating, the ductile fracture surface ratio at −20° C. is 50% or more. The thickness of the test piece is 0.6 to 3.0 mm, and when the thickness of the hot-rolled steel sheet exceeds 3.0 mm, the test piece taken from the hot-rolled steel sheet is ground on both sides to a thickness of 3.0 mm, and then subjected to the Charpy impact test.
In the present invention, "excellent resistance to delayed fracture after post-heating" means that no cracks are generated when a rectangular test piece having a sheared end surface, which is taken from the steel sheet after post-heating, is V-bent at 90°, the part opened by springback is tightened with a bolt or the like, and the piece is immersed in hydrochloric acid of pH 3 for 96 hours.
In the present invention, post-heating means a heat treatment in which the hot-rolled steel sheet (original sheet) is heated to 400° C. or higher.
また、本発明において、後加熱後の強度に優れるとは、後加熱後の熱延鋼板の強度の低下が、後加熱前の熱延鋼板（原板）の強度に対して、ビッカース硬さで５０以下であることを意味する。
本発明において、後加熱後の靭性に優れるとは、後加熱後の熱延鋼板から採取した試験片を用いたシャルピー衝撃試験において、－２０℃での延性破面率が５０％以上であることを意味する。なお、前記試験片の板厚は０．６～３．０ｍｍとし、熱延鋼板の板厚が３．０ｍｍを超える場合は、熱延鋼板から採取した試験片を板厚３．０ｍｍまで表裏面研削加工し、シャルピー衝撃試験に供するものとする。
本発明において、後加熱後の耐遅れ破壊特性に優れるとは、後加熱後の鋼板から採取した、せん断端面を有する短冊試験片に９０゜でＶ曲げ加工を施し、次いでスプリングバックで開いた分をボルトなどで締め付け、ｐＨ３の塩酸に９６ｈｒ浸漬した時に割れが発生しないことを意味する。
なお、本発明において、後加熱とは、熱延鋼板（原板）を４００℃以上に加熱する熱処理を意味する。 In the present invention, high strength means that the tensile strength (TS) is 1180 MPa or more and less than 1600 MPa.
In the present invention, excellent strength after post-heating means that the decrease in strength of the hot-rolled steel sheet after post-heating is 50 or less in Vickers hardness compared to the strength of the hot-rolled steel sheet (original sheet) before post-heating.
In the present invention, being excellent in toughness after post-heating means that in a Charpy impact test using a test piece taken from the hot-rolled steel sheet after post-heating, the ductile fracture surface ratio at −20° C. is 50% or more. The thickness of the test piece is 0.6 to 3.0 mm, and when the thickness of the hot-rolled steel sheet exceeds 3.0 mm, the test piece taken from the hot-rolled steel sheet is ground on both sides to a thickness of 3.0 mm, and then subjected to the Charpy impact test.
In the present invention, "excellent resistance to delayed fracture after post-heating" means that no cracks are generated when a rectangular test piece having a sheared end surface, which is taken from the steel sheet after post-heating, is V-bent at 90°, the part opened by springback is tightened with a bolt or the like, and the piece is immersed in hydrochloric acid of pH 3 for 96 hours.
In the present invention, post-heating means a heat treatment in which the hot-rolled steel sheet (original sheet) is heated to 400° C. or higher.
本発明は、以下の構成を有する。
［１］質量％で、
Ｃ：０．０６～０．２３％、
Ｓｉ：０．１～３．０％、
Ｍｎ：１．５～３．５％、
Ｐ：０％超０．０５０％以下、
Ｓ：０％超０．００５０％以下、
Ａｌ：０％超１．５％以下、
Ｎ：０％超０．０１０％以下、
Ｏ：０％超０．００３％以下を含み、さらに、
ＴｉとＮｂを合計で０．０４０～０．２００％含み、
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、
鋼組織は、マルテンサイトおよび／または下部ベイナイトを主相とし、残留オーステナイトが体積率で３％未満であり、
Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）が０．３００以上０．８００未満であり、
粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量が合計で０．０１０～０．０３０質量％であり、
表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度が１．８～５．０である、高強度熱延鋼板。
［２］前記成分組成が、さらに、質量％で、
Ｃｒ：０．００５～２．０％、
Ｎｉ：０．００５～２．０％、
Ｍｏ：０．００５～１．０％、
Ｖ：０．００５～０．５％、
Ｂ：０．０００２～０．００５０％、
Ｃａ：０．０００１～０．００５０％、
ＲＥＭ：０．０００１～０．００５０％
Ｃｕ：０．００５～０．５％、
Ｓｂ：０．００１０～０．１０％、および
Ｓｎ：０．００１０～０．１０％
のうちから選ばれる１種以上を含む、［１］に記載の高強度熱延鋼板。
［３］前記［１］または［２］に記載の高強度熱延鋼板の製造方法であって、
前記成分組成を有するスラブを１１５０～１３００℃の温度域に加熱し、該温度域で０．２～３．５時間保持し、
次いで、熱間圧延を施すに際し、
１０８０℃以上の温度域での合計圧下率が８０～９０％、９００℃以下の温度域での合計圧下率が２０％以上、かつ、下記式で求められるＴ（℃）以下での１パスあたりの圧下率が２５％以下となる条件で圧延した後、１．０ｓ以上放冷し、
次いで、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却し、５５０℃に達してから急冷を開始するまでの時間を０．５～４．０ｓとし、
次いで、１００～２５０℃の巻取り温度までを２００℃／ｓ以上の冷却速度で急冷し、前記巻取り温度で巻取る、高強度熱延鋼板の製造方法。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。 The present invention has the following configuration.
[1] In mass%,
C: 0.06 to 0.23%,
Si: 0.1 to 3.0%,
Mn: 1.5 to 3.5%,
P: more than 0% and not more than 0.050%;
S: more than 0% and 0.0050% or less;
Al: more than 0% and not more than 1.5%;
N: more than 0% and not more than 0.010%;
O: more than 0% and 0.003% or less;
Contains Ti and Nb in total 0.040 to 0.200%;
The balance is Fe and unavoidable impurities,
The steel structure has martensite and/or lower bainite as a main phase, and the volume fraction of retained austenite is less than 3%;
the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb, that is, (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb), is 0.300 or more and less than 0.800;
The total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is 0.010 to 0.030 mass %,
A high-strength hot-rolled steel sheet, in which the pole density of the {110}<111> orientation is 1.8 to 5.0 in a surface layer region extending from the surface to 100 μm in the sheet thickness center direction.
[2] The composition further comprises, in mass%,
Cr: 0.005 to 2.0%,
Ni: 0.005 to 2.0%,
Mo: 0.005 to 1.0%,
V: 0.005 to 0.5%,
B: 0.0002 to 0.0050%,
Ca: 0.0001 to 0.0050%,
REM: 0.0001 to 0.0050%
Cu: 0.005 to 0.5%,
Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%
The high strength hot rolled steel sheet according to [1], comprising one or more selected from the following:
[3] A method for producing a high strength hot rolled steel sheet according to the above [1] or [2],
A slab having the above-mentioned composition is heated to a temperature range of 1150 to 1300°C and held at that temperature range for 0.2 to 3.5 hours;
Next, when hot rolling is performed,
The total reduction in the temperature range of 1080°C or more is 80 to 90%, the total reduction in the temperature range of 900°C or less is 20% or more, and the reduction per pass at T (°C) or less calculated by the following formula is 25% or less. Then, the steel is allowed to cool for 1.0 s or more.
Next, the temperature is cooled at an average cooling rate of 50°C/s or more up to 550°C, and the time from reaching 550°C to starting quenching is set to 0.5 to 4.0 s.
Next, the steel sheet is quenched to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more, and then coiled at the coiling temperature.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
［１］質量％で、
Ｃ：０．０６～０．２３％、
Ｓｉ：０．１～３．０％、
Ｍｎ：１．５～３．５％、
Ｐ：０％超０．０５０％以下、
Ｓ：０％超０．００５０％以下、
Ａｌ：０％超１．５％以下、
Ｎ：０％超０．０１０％以下、
Ｏ：０％超０．００３％以下を含み、さらに、
ＴｉとＮｂを合計で０．０４０～０．２００％含み、
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、
鋼組織は、マルテンサイトおよび／または下部ベイナイトを主相とし、残留オーステナイトが体積率で３％未満であり、
Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）が０．３００以上０．８００未満であり、
粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量が合計で０．０１０～０．０３０質量％であり、
表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度が１．８～５．０である、高強度熱延鋼板。
［２］前記成分組成が、さらに、質量％で、
Ｃｒ：０．００５～２．０％、
Ｎｉ：０．００５～２．０％、
Ｍｏ：０．００５～１．０％、
Ｖ：０．００５～０．５％、
Ｂ：０．０００２～０．００５０％、
Ｃａ：０．０００１～０．００５０％、
ＲＥＭ：０．０００１～０．００５０％
Ｃｕ：０．００５～０．５％、
Ｓｂ：０．００１０～０．１０％、および
Ｓｎ：０．００１０～０．１０％
のうちから選ばれる１種以上を含む、［１］に記載の高強度熱延鋼板。
［３］前記［１］または［２］に記載の高強度熱延鋼板の製造方法であって、
前記成分組成を有するスラブを１１５０～１３００℃の温度域に加熱し、該温度域で０．２～３．５時間保持し、
次いで、熱間圧延を施すに際し、
１０８０℃以上の温度域での合計圧下率が８０～９０％、９００℃以下の温度域での合計圧下率が２０％以上、かつ、下記式で求められるＴ（℃）以下での１パスあたりの圧下率が２５％以下となる条件で圧延した後、１．０ｓ以上放冷し、
次いで、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却し、５５０℃に達してから急冷を開始するまでの時間を０．５～４．０ｓとし、
次いで、１００～２５０℃の巻取り温度までを２００℃／ｓ以上の冷却速度で急冷し、前記巻取り温度で巻取る、高強度熱延鋼板の製造方法。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。 The present invention has the following configuration.
[1] In mass%,
C: 0.06 to 0.23%,
Si: 0.1 to 3.0%,
Mn: 1.5 to 3.5%,
P: more than 0% and not more than 0.050%;
S: more than 0% and 0.0050% or less;
Al: more than 0% and not more than 1.5%;
N: more than 0% and not more than 0.010%;
O: more than 0% and 0.003% or less;
Contains Ti and Nb in total 0.040 to 0.200%;
The balance is Fe and unavoidable impurities,
The steel structure has martensite and/or lower bainite as a main phase, and the volume fraction of retained austenite is less than 3%;
the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb, that is, (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb), is 0.300 or more and less than 0.800;
The total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is 0.010 to 0.030 mass %,
A high-strength hot-rolled steel sheet, in which the pole density of the {110}<111> orientation is 1.8 to 5.0 in a surface layer region extending from the surface to 100 μm in the sheet thickness center direction.
[2] The composition further comprises, in mass%,
Cr: 0.005 to 2.0%,
Ni: 0.005 to 2.0%,
Mo: 0.005 to 1.0%,
V: 0.005 to 0.5%,
B: 0.0002 to 0.0050%,
Ca: 0.0001 to 0.0050%,
REM: 0.0001 to 0.0050%
Cu: 0.005 to 0.5%,
Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%
The high strength hot rolled steel sheet according to [1], comprising one or more selected from the following:
[3] A method for producing a high strength hot rolled steel sheet according to the above [1] or [2],
A slab having the above-mentioned composition is heated to a temperature range of 1150 to 1300°C and held at that temperature range for 0.2 to 3.5 hours;
Next, when hot rolling is performed,
The total reduction in the temperature range of 1080°C or more is 80 to 90%, the total reduction in the temperature range of 900°C or less is 20% or more, and the reduction per pass at T (°C) or less calculated by the following formula is 25% or less. Then, the steel is allowed to cool for 1.0 s or more.
Next, the temperature is cooled at an average cooling rate of 50°C/s or more up to 550°C, and the time from reaching 550°C to starting quenching is set to 0.5 to 4.0 s.
Next, the steel sheet is quenched to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more, and then coiled at the coiling temperature.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
本発明によれば、後加熱後の強度、靭性および耐遅れ破壊特性に優れる高強度熱延鋼板及びその製造方法を提供することができる。
The present invention provides a high-strength hot-rolled steel sheet that has excellent strength, toughness, and delayed fracture resistance after post-heating, and a manufacturing method thereof.
本発明によれば、自動車用部品の素材として好適な、後加熱後あるいは後加熱加工後の強度、靭性および耐遅れ破壊特性に優れる高強度熱延鋼板が得られる。
本発明の高強度熱延鋼板を用いれば、加工性改善や疲労特性向上等のために加熱処理を施した後でも、高強度と良好な靭性および優れた耐遅れ破壊特性を発現する高強度自動車部品等の製品を得ることができる。 According to the present invention, it is possible to obtain a high-strength hot-rolled steel sheet which is suitable as a material for automobile parts and which has excellent strength, toughness and delayed fracture resistance after post-heating or after post-heat processing.
By using the high-strength hot-rolled steel sheet of the present invention, it is possible to obtain products such as high-strength automobile parts that exhibit high strength, good toughness, and excellent delayed fracture resistance even after heat treatment is performed to improve workability and fatigue properties.
本発明の高強度熱延鋼板を用いれば、加工性改善や疲労特性向上等のために加熱処理を施した後でも、高強度と良好な靭性および優れた耐遅れ破壊特性を発現する高強度自動車部品等の製品を得ることができる。 According to the present invention, it is possible to obtain a high-strength hot-rolled steel sheet which is suitable as a material for automobile parts and which has excellent strength, toughness and delayed fracture resistance after post-heating or after post-heat processing.
By using the high-strength hot-rolled steel sheet of the present invention, it is possible to obtain products such as high-strength automobile parts that exhibit high strength, good toughness, and excellent delayed fracture resistance even after heat treatment is performed to improve workability and fatigue properties.
以下に、本発明の高強度熱延鋼板及びその製造方法の実施形態について詳細に説明する。なお、本発明は、以下の実施形態に限定されない。
Below, an embodiment of the high-strength hot-rolled steel sheet and its manufacturing method of the present invention will be described in detail. Note that the present invention is not limited to the following embodiment.
＜高強度熱延鋼板＞
本発明の高強度熱延鋼板は、熱間圧延ままの黒皮、または熱間圧延後さらに酸洗する白皮と称される熱延鋼板のどちらであってもよい。また、本発明が目的とする高強度熱延鋼板は、板厚が０．６ｍｍ以上であることが好ましい。また、本発明の高強度熱延鋼板は、板厚が１０．０ｍｍ以下であることが好ましい。本発明の高強度熱延鋼板を自動車用部品の素材として用いる場合には、板厚が１．０ｍｍ以上であることがより好ましい。また、本発明の高強度熱延鋼板を自動車用部品の素材として用いる場合には、板厚が６．０ｍｍ以下であることがより好ましい。また、本発明の高強度熱延鋼板の板幅は、５００ｍｍ以上であることが好ましく、７００ｍｍ以上であることがより好ましい。本発明の高強度熱延鋼板の板幅は、１８００ｍｍ以下であることが好ましく、１４００ｍｍ以下であることがより好ましい。 <High-strength hot-rolled steel sheet>
The high-strength hot-rolled steel sheet of the present invention may be either a black skin as hot-rolled, or a hot-rolled steel sheet called a white skin which is further pickled after hot rolling. The high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 0.6 mm or more. The high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 10.0 mm or less. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 1.0 mm or more. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 6.0 mm or less. The width of the high-strength hot-rolled steel sheet of the present invention is preferably 500 mm or more, more preferably 700 mm or more. The width of the high-strength hot-rolled steel sheet of the present invention is preferably 1800 mm or less, more preferably 1400 mm or less.
本発明の高強度熱延鋼板は、熱間圧延ままの黒皮、または熱間圧延後さらに酸洗する白皮と称される熱延鋼板のどちらであってもよい。また、本発明が目的とする高強度熱延鋼板は、板厚が０．６ｍｍ以上であることが好ましい。また、本発明の高強度熱延鋼板は、板厚が１０．０ｍｍ以下であることが好ましい。本発明の高強度熱延鋼板を自動車用部品の素材として用いる場合には、板厚が１．０ｍｍ以上であることがより好ましい。また、本発明の高強度熱延鋼板を自動車用部品の素材として用いる場合には、板厚が６．０ｍｍ以下であることがより好ましい。また、本発明の高強度熱延鋼板の板幅は、５００ｍｍ以上であることが好ましく、７００ｍｍ以上であることがより好ましい。本発明の高強度熱延鋼板の板幅は、１８００ｍｍ以下であることが好ましく、１４００ｍｍ以下であることがより好ましい。 <High-strength hot-rolled steel sheet>
The high-strength hot-rolled steel sheet of the present invention may be either a black skin as hot-rolled, or a hot-rolled steel sheet called a white skin which is further pickled after hot rolling. The high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 0.6 mm or more. The high-strength hot-rolled steel sheet of the present invention preferably has a thickness of 10.0 mm or less. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 1.0 mm or more. When the high-strength hot-rolled steel sheet of the present invention is used as a material for automobile parts, the thickness is more preferably 6.0 mm or less. The width of the high-strength hot-rolled steel sheet of the present invention is preferably 500 mm or more, more preferably 700 mm or more. The width of the high-strength hot-rolled steel sheet of the present invention is preferably 1800 mm or less, more preferably 1400 mm or less.
本発明の高強度熱延鋼板は、特定の成分組成と、特定の鋼組織とを有する。ここでは、成分組成、鋼組織の順に説明する。
The high-strength hot-rolled steel sheet of the present invention has a specific chemical composition and a specific steel structure. Here, the chemical composition and steel structure will be explained in that order.
まず、本発明の高強度熱延鋼板の成分組成について説明する。なお、成分組成の含有量を表す「％」は「質量％」を意味するものとする。
First, we will explain the composition of the high-strength hot-rolled steel sheet of the present invention. Note that "%" representing the content of the composition means "mass %."
本発明の高強度熱延鋼板の成分組成は、質量％で、Ｃ：０．０６～０．２３％、Ｓｉ：０．１～３．０％、Ｍｎ：１．５～３．５％、Ｐ：０．０５０％以下（０％を含まない）、Ｓ：０．００５０％以下（０％を含まない）、Ａｌ：１．５％以下（０％を含まない）、Ｎ：０．０１０％以下（０％を含まない）、Ｏ：０．００３％以下（０％を含まない）、ＴｉとＮｂを合計で０．０４０～０．２００％含み、残部がＦｅおよび不可避的不純物からなる。
The composition of the high-strength hot-rolled steel sheet of the present invention is, in mass%, C: 0.06-0.23%, Si: 0.1-3.0%, Mn: 1.5-3.5%, P: 0.050% or less (excluding 0%), S: 0.0050% or less (excluding 0%), Al: 1.5% or less (excluding 0%), N: 0.010% or less (excluding 0%), O: 0.003% or less (excluding 0%), Ti and Nb in total 0.040-0.200%, and the balance consisting of Fe and unavoidable impurities.
Ｃ：０．０６～０．２３％
Ｃは、マルテンサイトや下部ベイナイトを生成および強化させてＴＳを上昇させたり、Ｔｉ、Ｎｂ、Ｎ等と結合して析出物を生成することで後加熱後の強度低下抑制等に有効な元素である。Ｃ含有量が０．０６％未満ではこのような効果が十分得られず、１１８０ＭＰａ以上の鋼板（原板）のＴＳあるいは後加熱後の優れた強度が得られない。一方、Ｃ含有量が０．２３％を超えると後加熱後の靭性の低下が顕著となり、後加熱後の優れた靭性が得られない。したがって、Ｃ含有量は０．０６～０．２３％とする。Ｃ含有量は、好ましくは０．０７％以上とする。また、Ｃ含有量は、好ましくは０．２２％以下とし、より好ましくは０．２０％以下とする。 C: 0.06 to 0.23%
C is an element that is effective in increasing TS by generating and strengthening martensite and lower bainite, and in suppressing strength reduction after post-heating by combining with Ti, Nb, N, etc. to generate precipitates. If the C content is less than 0.06%, such effects cannot be sufficiently obtained, and the TS of the steel plate (original plate) of 1180 MPa or more or excellent strength after post-heating cannot be obtained. On the other hand, if the C content exceeds 0.23%, the decrease in toughness after post-heating becomes significant, and excellent toughness after post-heating cannot be obtained. Therefore, the C content is set to 0.06 to 0.23%. The C content is preferably set to 0.07% or more. The C content is preferably set to 0.22% or less, and more preferably set to 0.20% or less.
Ｃは、マルテンサイトや下部ベイナイトを生成および強化させてＴＳを上昇させたり、Ｔｉ、Ｎｂ、Ｎ等と結合して析出物を生成することで後加熱後の強度低下抑制等に有効な元素である。Ｃ含有量が０．０６％未満ではこのような効果が十分得られず、１１８０ＭＰａ以上の鋼板（原板）のＴＳあるいは後加熱後の優れた強度が得られない。一方、Ｃ含有量が０．２３％を超えると後加熱後の靭性の低下が顕著となり、後加熱後の優れた靭性が得られない。したがって、Ｃ含有量は０．０６～０．２３％とする。Ｃ含有量は、好ましくは０．０７％以上とする。また、Ｃ含有量は、好ましくは０．２２％以下とし、より好ましくは０．２０％以下とする。 C: 0.06 to 0.23%
C is an element that is effective in increasing TS by generating and strengthening martensite and lower bainite, and in suppressing strength reduction after post-heating by combining with Ti, Nb, N, etc. to generate precipitates. If the C content is less than 0.06%, such effects cannot be sufficiently obtained, and the TS of the steel plate (original plate) of 1180 MPa or more or excellent strength after post-heating cannot be obtained. On the other hand, if the C content exceeds 0.23%, the decrease in toughness after post-heating becomes significant, and excellent toughness after post-heating cannot be obtained. Therefore, the C content is set to 0.06 to 0.23%. The C content is preferably set to 0.07% or more. The C content is preferably set to 0.22% or less, and more preferably set to 0.20% or less.
Ｓｉ：０．１～３．０％
Ｓｉは、鋼の固溶強化および後加熱後の強度低下の抑制に有効な元素である。このような効果を得るには、Ｓｉ含有量を０．１％以上とする必要がある。一方、Ｓｉ含有量が３．０％を超えると、ポリゴナルフェライトが過剰に生成して本発明の鋼組織が得られなくなる。したがって、Ｓｉ含有量は０．１～３．０％とする。Ｓｉ含有量は、好ましくは０．２％以上とする。また、Ｓｉ含有量は、好ましくは２．０％以下とし、より好ましくは１．５％以下とする。 Si: 0.1 to 3.0%
Silicon is an element effective in solution strengthening of steel and suppressing the decrease in strength after post-heating. To obtain such an effect, the silicon content must be 0.1% or more. On the other hand, if the silicon content exceeds 3.0%, polygonal ferrite is excessively formed and the steel structure of the present invention cannot be obtained. Therefore, the silicon content is set to 0.1 to 3.0%. The silicon content is preferably set to 0.2% or more. The silicon content is preferably set to 2.0% or less, more preferably 1.5% or less.
Ｓｉは、鋼の固溶強化および後加熱後の強度低下の抑制に有効な元素である。このような効果を得るには、Ｓｉ含有量を０．１％以上とする必要がある。一方、Ｓｉ含有量が３．０％を超えると、ポリゴナルフェライトが過剰に生成して本発明の鋼組織が得られなくなる。したがって、Ｓｉ含有量は０．１～３．０％とする。Ｓｉ含有量は、好ましくは０．２％以上とする。また、Ｓｉ含有量は、好ましくは２．０％以下とし、より好ましくは１．５％以下とする。 Si: 0.1 to 3.0%
Silicon is an element effective in solution strengthening of steel and suppressing the decrease in strength after post-heating. To obtain such an effect, the silicon content must be 0.1% or more. On the other hand, if the silicon content exceeds 3.0%, polygonal ferrite is excessively formed and the steel structure of the present invention cannot be obtained. Therefore, the silicon content is set to 0.1 to 3.0%. The silicon content is preferably set to 0.2% or more. The silicon content is preferably set to 2.0% or less, more preferably 1.5% or less.
Ｍｎ：１．５～３．５％
Ｍｎは、フェライトや上部ベイナイトを抑制して、下部ベイナイトやマルテンサイトを生成させるのに有効な元素である。Ｍｎ含有量が１．５％未満ではこうした効果が十分得られず、ポリゴナルフェライトや上部ベイナイト等が生成し、本発明のミクロ組織が得られなくなる。一方、Ｍｎ含有量が３．５％を超えると、靭性の低下や耐遅れ破壊特性の低下が顕著になり、後加熱後の優れた靭性と耐遅れ破壊特性が得られなくなる。したがって、Ｍｎ含有量は１．５～３．５％とする。Ｍｎ含有量は、好ましくは１．６％以上とする。また、Ｍｎ含有量は、好ましくは３．０％以下とし、より好ましくは２．５％以下とする。 Mn: 1.5 to 3.5%
Mn is an element effective in suppressing ferrite and upper bainite and generating lower bainite and martensite. If the Mn content is less than 1.5%, this effect is not sufficiently obtained, and polygonal ferrite, upper bainite, etc. are generated, and the microstructure of the present invention cannot be obtained. On the other hand, if the Mn content exceeds 3.5%, the deterioration of toughness and delayed fracture resistance becomes significant, and excellent toughness and delayed fracture resistance after post-heating cannot be obtained. Therefore, the Mn content is set to 1.5 to 3.5%. The Mn content is preferably 1.6% or more. The Mn content is preferably 3.0% or less, more preferably 2.5% or less.
Ｍｎは、フェライトや上部ベイナイトを抑制して、下部ベイナイトやマルテンサイトを生成させるのに有効な元素である。Ｍｎ含有量が１．５％未満ではこうした効果が十分得られず、ポリゴナルフェライトや上部ベイナイト等が生成し、本発明のミクロ組織が得られなくなる。一方、Ｍｎ含有量が３．５％を超えると、靭性の低下や耐遅れ破壊特性の低下が顕著になり、後加熱後の優れた靭性と耐遅れ破壊特性が得られなくなる。したがって、Ｍｎ含有量は１．５～３．５％とする。Ｍｎ含有量は、好ましくは１．６％以上とする。また、Ｍｎ含有量は、好ましくは３．０％以下とし、より好ましくは２．５％以下とする。 Mn: 1.5 to 3.5%
Mn is an element effective in suppressing ferrite and upper bainite and generating lower bainite and martensite. If the Mn content is less than 1.5%, this effect is not sufficiently obtained, and polygonal ferrite, upper bainite, etc. are generated, and the microstructure of the present invention cannot be obtained. On the other hand, if the Mn content exceeds 3.5%, the deterioration of toughness and delayed fracture resistance becomes significant, and excellent toughness and delayed fracture resistance after post-heating cannot be obtained. Therefore, the Mn content is set to 1.5 to 3.5%. The Mn content is preferably 1.6% or more. The Mn content is preferably 3.0% or less, more preferably 2.5% or less.
Ｐ：０％超０．０５０％以下
Ｐは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが望ましい。本発明ではＰ含有量が０．０５０％まで許容できる。したがって、Ｐ含有量は０．０５０％以下とする。Ｐ含有量は、好ましくは０．０３０％以下とする。下限は特に限定されず、Ｐ含有量は０％超でよいが、Ｐ含有量が０．００１％未満では生産能率が低下するため、Ｐ含有量は０．００１％以上が好ましい。 P: more than 0% and 0.050% or less P reduces the toughness and delayed fracture resistance after post-heating, so it is desirable to reduce the amount as much as possible. In the present invention, a P content of up to 0.050% is acceptable. Therefore, the P content is set to 0.050% or less. The P content is preferably set to 0.030% or less. There is no particular lower limit, and the P content may be more than 0%, but if the P content is less than 0.001%, the production efficiency decreases, so the P content is preferably 0.001% or more.
Ｐは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが望ましい。本発明ではＰ含有量が０．０５０％まで許容できる。したがって、Ｐ含有量は０．０５０％以下とする。Ｐ含有量は、好ましくは０．０３０％以下とする。下限は特に限定されず、Ｐ含有量は０％超でよいが、Ｐ含有量が０．００１％未満では生産能率が低下するため、Ｐ含有量は０．００１％以上が好ましい。 P: more than 0% and 0.050% or less P reduces the toughness and delayed fracture resistance after post-heating, so it is desirable to reduce the amount as much as possible. In the present invention, a P content of up to 0.050% is acceptable. Therefore, the P content is set to 0.050% or less. The P content is preferably set to 0.030% or less. There is no particular lower limit, and the P content may be more than 0%, but if the P content is less than 0.001%, the production efficiency decreases, so the P content is preferably 0.001% or more.
Ｓ：０％超０．００５０％以下
Ｓは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが好ましい。本発明ではＳ含有量が０．００５０％まで許容できる。したがって、Ｓ含有量は０．００５０％以下とする。Ｓ含有量は、好ましくは０．００３０％以下とし、より好ましくは０．００２０％以下とし、さらに好ましくは０．００１５％とする。下限は特に限定されず、Ｓ含有量は０％超でよいが、Ｓ含有量が０．０００２％未満では生産能率が低下するため、Ｓ含有量は０．０００２％以上が好ましい。 S: more than 0% and 0.0050% or less S reduces the toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible. In the present invention, the S content can be up to 0.0050%. Therefore, the S content is set to 0.0050% or less. The S content is preferably set to 0.0030% or less, more preferably set to 0.0020% or less, and even more preferably set to 0.0015%. There is no particular lower limit, and the S content may be more than 0%, but if the S content is less than 0.0002%, the production efficiency decreases, so the S content is preferably 0.0002% or more.
Ｓは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが好ましい。本発明ではＳ含有量が０．００５０％まで許容できる。したがって、Ｓ含有量は０．００５０％以下とする。Ｓ含有量は、好ましくは０．００３０％以下とし、より好ましくは０．００２０％以下とし、さらに好ましくは０．００１５％とする。下限は特に限定されず、Ｓ含有量は０％超でよいが、Ｓ含有量が０．０００２％未満では生産能率が低下するため、Ｓ含有量は０．０００２％以上が好ましい。 S: more than 0% and 0.0050% or less S reduces the toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible. In the present invention, the S content can be up to 0.0050%. Therefore, the S content is set to 0.0050% or less. The S content is preferably set to 0.0030% or less, more preferably set to 0.0020% or less, and even more preferably set to 0.0015%. There is no particular lower limit, and the S content may be more than 0%, but if the S content is less than 0.0002%, the production efficiency decreases, so the S content is preferably 0.0002% or more.
Ａｌ：０％超１．５％以下
Ａｌは、脱酸剤として作用し、脱酸工程で添加することが好ましい。Ａｌ含有量は０％超でよいが、脱酸剤として用いる観点からは、Ａｌ含有量は０．０１％以上が好ましい。一方、多量にＡｌを含有するとポリゴナルフェライトが多量に生成して本発明の鋼組織が得られなくなる。本発明ではＡｌ含有量が１．５％まで許容される。したがって、Ａｌ含有量は１．５％以下とする。Ａｌ含有量は、好ましくは０．５０％以下、より好ましくは０．２０％以下とする。 Al: more than 0% and not more than 1.5% Al acts as a deoxidizer, and is preferably added in the deoxidization process. The Al content may be more than 0%, but from the viewpoint of using it as a deoxidizer, the Al content is preferably 0.01% or more. On the other hand, if a large amount of Al is contained, a large amount of polygonal ferrite is generated, and the steel structure of the present invention cannot be obtained. In the present invention, the Al content is allowed up to 1.5%. Therefore, the Al content is set to 1.5% or less. The Al content is preferably set to 0.50% or less, more preferably 0.20% or less.
Ａｌは、脱酸剤として作用し、脱酸工程で添加することが好ましい。Ａｌ含有量は０％超でよいが、脱酸剤として用いる観点からは、Ａｌ含有量は０．０１％以上が好ましい。一方、多量にＡｌを含有するとポリゴナルフェライトが多量に生成して本発明の鋼組織が得られなくなる。本発明ではＡｌ含有量が１．５％まで許容される。したがって、Ａｌ含有量は１．５％以下とする。Ａｌ含有量は、好ましくは０．５０％以下、より好ましくは０．２０％以下とする。 Al: more than 0% and not more than 1.5% Al acts as a deoxidizer, and is preferably added in the deoxidization process. The Al content may be more than 0%, but from the viewpoint of using it as a deoxidizer, the Al content is preferably 0.01% or more. On the other hand, if a large amount of Al is contained, a large amount of polygonal ferrite is generated, and the steel structure of the present invention cannot be obtained. In the present invention, the Al content is allowed up to 1.5%. Therefore, the Al content is set to 1.5% or less. The Al content is preferably set to 0.50% or less, more preferably 0.20% or less.
Ｎ：０％超０．０１０％以下
Ｎは、ＴｉＮやＮｂＣを生成させ、微細なＴｉＣやＮｂＣ等の析出を阻害するため、その量は極力低減することが好ましい。本発明ではＮ含有量が０．０１０％まで許容できる。したがって、Ｎ含有量は０．０１０％以下とする。Ｎ含有量は、好ましくは０．００７％以下とする。下限は特に限定されず、Ｎ含有量は０％超でよいが、Ｎ含有量が０．０００５％未満では生産能率が低下するため、Ｎ含有量は０．０００５％以上が好ましい。 N: more than 0% and 0.010% or less N generates TiN and NbC and inhibits the precipitation of fine TiC, NbC, etc., so it is preferable to reduce the amount as much as possible. In the present invention, an N content of up to 0.010% is acceptable. Therefore, the N content is set to 0.010% or less. The N content is preferably set to 0.007% or less. There is no particular lower limit, and the N content may be more than 0%, but if the N content is less than 0.0005%, the production efficiency decreases, so the N content is preferably 0.0005% or more.
Ｎは、ＴｉＮやＮｂＣを生成させ、微細なＴｉＣやＮｂＣ等の析出を阻害するため、その量は極力低減することが好ましい。本発明ではＮ含有量が０．０１０％まで許容できる。したがって、Ｎ含有量は０．０１０％以下とする。Ｎ含有量は、好ましくは０．００７％以下とする。下限は特に限定されず、Ｎ含有量は０％超でよいが、Ｎ含有量が０．０００５％未満では生産能率が低下するため、Ｎ含有量は０．０００５％以上が好ましい。 N: more than 0% and 0.010% or less N generates TiN and NbC and inhibits the precipitation of fine TiC, NbC, etc., so it is preferable to reduce the amount as much as possible. In the present invention, an N content of up to 0.010% is acceptable. Therefore, the N content is set to 0.010% or less. The N content is preferably set to 0.007% or less. There is no particular lower limit, and the N content may be more than 0%, but if the N content is less than 0.0005%, the production efficiency decreases, so the N content is preferably 0.0005% or more.
Ｏ：０％超０．００３％以下
Ｏは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが好ましい。本発明ではＯ含有量が０．００３％まで許容できる。したがって、Ｏ含有量は０．００３％以下とする。Ｏ含有量は、好ましくは０．００２％以下とする。下限は特に限定されず、Ｏ含有量は０％超でよいが、Ｏ含有量が０．０００２％未満では生産能率が低下するため、Ｏ含有量は０．０００２％以上が好ましい。 O: more than 0% and 0.003% or less O reduces toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible. In the present invention, an O content of up to 0.003% is acceptable. Therefore, the O content is set to 0.003% or less. The O content is preferably set to 0.002% or less. There is no particular lower limit, and the O content may be more than 0%, but if the O content is less than 0.0002%, production efficiency decreases, so the O content is preferably 0.0002% or more.
Ｏは、後加熱後の靭性や耐遅れ破壊特性を低下させるため、その量は極力低減することが好ましい。本発明ではＯ含有量が０．００３％まで許容できる。したがって、Ｏ含有量は０．００３％以下とする。Ｏ含有量は、好ましくは０．００２％以下とする。下限は特に限定されず、Ｏ含有量は０％超でよいが、Ｏ含有量が０．０００２％未満では生産能率が低下するため、Ｏ含有量は０．０００２％以上が好ましい。 O: more than 0% and 0.003% or less O reduces toughness and delayed fracture resistance after post-heating, so it is preferable to reduce the amount as much as possible. In the present invention, an O content of up to 0.003% is acceptable. Therefore, the O content is set to 0.003% or less. The O content is preferably set to 0.002% or less. There is no particular lower limit, and the O content may be more than 0%, but if the O content is less than 0.0002%, production efficiency decreases, so the O content is preferably 0.0002% or more.
ＴｉとＮｂの合計：０．０４０～０．２００％
ＴｉおよびＮｂは、本発明において最も重要な元素であり、後加熱後に適度なＴｉＣやＮｂＣ等の微細析出物を生成させることで、後加熱後の優れた強度、靭性および耐遅れ破壊特性を得るのに必要な元素である。ＴｉとＮｂの含有量の合計が０．０４０％未満ではこのような効果が十分得られず、後加熱後の優れた強度が得られない。一方、ＴｉとＮｂの含有量の合計が０．２００％を超えると粗大なＴｉやＮｂを含有する析出物が多くなることで後加熱後の耐遅れ破壊特性の低下を招くほか、後加熱後の析出物が過剰となり、後加熱後の優れた靭性が得られなくなる。したがって、ＴｉとＮｂの含有量は合計で０．０４０～０．２００％とする。ＴｉとＮｂの含有量は合計で、好ましくは０．０５０％以上とし、より好ましくは０．０６０％以上とする。また、ＴｉとＮｂの含有量は合計で、好ましくは０．１６０％以下とし、より好ましくは０．１２０％以下とする。なお、ＴｉとＮｂの含有量は、合計で上記範囲を満たせばよく、どちらか一方の含有量は０％であってもよい。 Sum of Ti and Nb: 0.040 to 0.200%
Ti and Nb are the most important elements in the present invention, and are necessary elements for obtaining excellent strength, toughness, and delayed fracture resistance properties after post-heating by generating appropriate fine precipitates such as TiC and NbC after post-heating. If the total content of Ti and Nb is less than 0.040%, such effects are not sufficiently obtained, and excellent strength after post-heating is not obtained. On the other hand, if the total content of Ti and Nb exceeds 0.200%, the amount of coarse precipitates containing Ti and Nb increases, which leads to a decrease in delayed fracture resistance properties after post-heating, and the precipitates after post-heating become excessive, making it impossible to obtain excellent toughness after post-heating. Therefore, the total content of Ti and Nb is set to 0.040 to 0.200%. The total content of Ti and Nb is preferably 0.050% or more, more preferably 0.060% or more. The total content of Ti and Nb is preferably 0.160% or less, more preferably 0.120% or less. The total content of Ti and Nb needs to be within the above range, and the content of either one of them may be 0%.
ＴｉおよびＮｂは、本発明において最も重要な元素であり、後加熱後に適度なＴｉＣやＮｂＣ等の微細析出物を生成させることで、後加熱後の優れた強度、靭性および耐遅れ破壊特性を得るのに必要な元素である。ＴｉとＮｂの含有量の合計が０．０４０％未満ではこのような効果が十分得られず、後加熱後の優れた強度が得られない。一方、ＴｉとＮｂの含有量の合計が０．２００％を超えると粗大なＴｉやＮｂを含有する析出物が多くなることで後加熱後の耐遅れ破壊特性の低下を招くほか、後加熱後の析出物が過剰となり、後加熱後の優れた靭性が得られなくなる。したがって、ＴｉとＮｂの含有量は合計で０．０４０～０．２００％とする。ＴｉとＮｂの含有量は合計で、好ましくは０．０５０％以上とし、より好ましくは０．０６０％以上とする。また、ＴｉとＮｂの含有量は合計で、好ましくは０．１６０％以下とし、より好ましくは０．１２０％以下とする。なお、ＴｉとＮｂの含有量は、合計で上記範囲を満たせばよく、どちらか一方の含有量は０％であってもよい。 Sum of Ti and Nb: 0.040 to 0.200%
Ti and Nb are the most important elements in the present invention, and are necessary elements for obtaining excellent strength, toughness, and delayed fracture resistance properties after post-heating by generating appropriate fine precipitates such as TiC and NbC after post-heating. If the total content of Ti and Nb is less than 0.040%, such effects are not sufficiently obtained, and excellent strength after post-heating is not obtained. On the other hand, if the total content of Ti and Nb exceeds 0.200%, the amount of coarse precipitates containing Ti and Nb increases, which leads to a decrease in delayed fracture resistance properties after post-heating, and the precipitates after post-heating become excessive, making it impossible to obtain excellent toughness after post-heating. Therefore, the total content of Ti and Nb is set to 0.040 to 0.200%. The total content of Ti and Nb is preferably 0.050% or more, more preferably 0.060% or more. The total content of Ti and Nb is preferably 0.160% or less, more preferably 0.120% or less. The total content of Ti and Nb needs to be within the above range, and the content of either one of them may be 0%.
上記成分が本発明の高強度熱延鋼板の基本の成分である。本発明の高強度熱延鋼板は、上記成分を含有し、残部はＦｅおよび不可避的不純物の成分組成とすることができる。
The above components are the basic components of the high-strength hot-rolled steel sheet of the present invention. The high-strength hot-rolled steel sheet of the present invention contains the above components, with the remainder being Fe and unavoidable impurities.
本発明の高強度熱延鋼板は、上記成分に加えて、さらに、Ｃｒ：０．００５～２．０％、Ｎｉ：０．００５～２．０％、Ｍｏ：０．００５～１．０％、Ｖ：０．００５～０．５％、Ｂ：０．０００２～０．００５０％、Ｃａ：０．０００１～０．００５０％、ＲＥＭ：０．０００１～０．００５０％、Ｃｕ：０．００５～０．５％、Ｓｂ：０．００１０～０．１０％、Ｓｎ：０．００１０～０．１０％のうちから選ばれる１種以上を含有することができる。
In addition to the above components, the high-strength hot-rolled steel sheet of the present invention can further contain one or more selected from Cr: 0.005-2.0%, Ni: 0.005-2.0%, Mo: 0.005-1.0%, V: 0.005-0.5%, B: 0.0002-0.0050%, Ca: 0.0001-0.0050%, REM: 0.0001-0.0050%, Cu: 0.005-0.5%, Sb: 0.0010-0.10%, Sn: 0.0010-0.10%.
Ｃｒ：０．００５～２．０％
Ｃｒは、フェライトを抑制して、下部ベイナイトやマルテンサイトを生成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、Ｃｒを含有する場合、Ｃｒ含有量を０．００５％以上とすることが好ましい。一方、Ｃｒの含有量が２．０％を超えると、耐食性の低下が顕著となる場合があるため、Ｃｒを含有する場合、Ｃｒ含有量を２．０％以下とすることが好ましい。Ｃｒ含有量は、より好ましくは０．１％以上とする。また、Ｃｒ含有量は、より好ましくは０．８％以下とする。 Cr: 0.005 to 2.0%
Cr is an element effective in suppressing ferrite and generating lower bainite and martensite. In order to obtain such an effect, when Cr is contained, the Cr content is preferably 0.005% or more. On the other hand, when the Cr content exceeds 2.0%, the corrosion resistance may be significantly decreased, so when Cr is contained, the Cr content is preferably 2.0% or less. The Cr content is more preferably 0.1% or more. Moreover, the Cr content is more preferably 0.8% or less.
Ｃｒは、フェライトを抑制して、下部ベイナイトやマルテンサイトを生成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、Ｃｒを含有する場合、Ｃｒ含有量を０．００５％以上とすることが好ましい。一方、Ｃｒの含有量が２．０％を超えると、耐食性の低下が顕著となる場合があるため、Ｃｒを含有する場合、Ｃｒ含有量を２．０％以下とすることが好ましい。Ｃｒ含有量は、より好ましくは０．１％以上とする。また、Ｃｒ含有量は、より好ましくは０．８％以下とする。 Cr: 0.005 to 2.0%
Cr is an element effective in suppressing ferrite and generating lower bainite and martensite. In order to obtain such an effect, when Cr is contained, the Cr content is preferably 0.005% or more. On the other hand, when the Cr content exceeds 2.0%, the corrosion resistance may be significantly decreased, so when Cr is contained, the Cr content is preferably 2.0% or less. The Cr content is more preferably 0.1% or more. Moreover, the Cr content is more preferably 0.8% or less.
Ｎｉ：０．００５～２．０％
Ｎｉは、フェライトを抑制して、下部ベイナイトやマルテンサイトを生成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、Ｎｉを含有する場合、Ｎｉ含有量を０．００５％以上とすることが好ましい。一方、Ｎｉの含有量が２．０％を超えると、残留γが多量に形成して後加熱後の靭性低下を招く場合があるため、Ｎｉを含有する場合、Ｎｉ含有量を２．０％以下とすることが好ましい。Ｎｉ含有量は、より好ましくは０．０５％以上とする。また、Ｎｉ含有量は、より好ましくは０．８％以下とし、さらに好ましくは０．５％以下とする。 Ni: 0.005 to 2.0%
Ni is an element effective in suppressing ferrite and generating lower bainite and martensite. In order to obtain such an effect, when Ni is contained, the Ni content is preferably 0.005% or more. On the other hand, when the Ni content exceeds 2.0%, a large amount of residual γ is formed, which may lead to a decrease in toughness after post-heating, so when Ni is contained, the Ni content is preferably 2.0% or less. The Ni content is more preferably 0.05% or more. Moreover, the Ni content is more preferably 0.8% or less, and further preferably 0.5% or less.
Ｎｉは、フェライトを抑制して、下部ベイナイトやマルテンサイトを生成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、Ｎｉを含有する場合、Ｎｉ含有量を０．００５％以上とすることが好ましい。一方、Ｎｉの含有量が２．０％を超えると、残留γが多量に形成して後加熱後の靭性低下を招く場合があるため、Ｎｉを含有する場合、Ｎｉ含有量を２．０％以下とすることが好ましい。Ｎｉ含有量は、より好ましくは０．０５％以上とする。また、Ｎｉ含有量は、より好ましくは０．８％以下とし、さらに好ましくは０．５％以下とする。 Ni: 0.005 to 2.0%
Ni is an element effective in suppressing ferrite and generating lower bainite and martensite. In order to obtain such an effect, when Ni is contained, the Ni content is preferably 0.005% or more. On the other hand, when the Ni content exceeds 2.0%, a large amount of residual γ is formed, which may lead to a decrease in toughness after post-heating, so when Ni is contained, the Ni content is preferably 2.0% or less. The Ni content is more preferably 0.05% or more. Moreover, the Ni content is more preferably 0.8% or less, and further preferably 0.5% or less.
Ｍｏ：０．００５～１．０％
Ｍｏは、鋼板の焼き入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｍｏを含有する場合、Ｍｏ含有量を０．００５％以上とすることが好ましい。一方、Ｍｏの含有量が１．０％を超えると、Ｍｏ系析出物の生成が顕著となり、後加熱後の靭性低下を招く場合があるため、Ｍｏを含有する場合、Ｍｏ含有量を１．０％以下とすることが好ましい。Ｍｏ含有量は、より好ましくは０．０５％以上とする。また、Ｍｏ含有量は、より好ましくは０．５０％以下とする。 Mo: 0.005 to 1.0%
Mo is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when Mo is contained, the Mo content is preferably 0.005% or more. On the other hand, when the Mo content exceeds 1.0%, the generation of Mo-based precipitates becomes significant, which may lead to a decrease in toughness after post-heating, so when Mo is contained, the Mo content is preferably 1.0% or less. The Mo content is more preferably 0.05% or more. Moreover, the Mo content is more preferably 0.50% or less.
Ｍｏは、鋼板の焼き入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｍｏを含有する場合、Ｍｏ含有量を０．００５％以上とすることが好ましい。一方、Ｍｏの含有量が１．０％を超えると、Ｍｏ系析出物の生成が顕著となり、後加熱後の靭性低下を招く場合があるため、Ｍｏを含有する場合、Ｍｏ含有量を１．０％以下とすることが好ましい。Ｍｏ含有量は、より好ましくは０．０５％以上とする。また、Ｍｏ含有量は、より好ましくは０．５０％以下とする。 Mo: 0.005 to 1.0%
Mo is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when Mo is contained, the Mo content is preferably 0.005% or more. On the other hand, when the Mo content exceeds 1.0%, the generation of Mo-based precipitates becomes significant, which may lead to a decrease in toughness after post-heating, so when Mo is contained, the Mo content is preferably 1.0% or less. The Mo content is more preferably 0.05% or more. Moreover, the Mo content is more preferably 0.50% or less.
Ｖ：０．００５～０．５％
Ｖは、鋼板の焼き入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｖを含有する場合、Ｖ含有量を０．００５％以上とすることが好ましい。一方、Ｖの含有量が０．５％を超えると、Ｖ系析出物の生成が過剰となり、後加熱後の靭性低下を招く場合があるため、Ｖを含有する場合、Ｖ含有量を０．５％以下とすることが好ましい。Ｖ含有量は、より好ましくは０．０１％以上とする。また、Ｖ含有量は、より好ましくは０．１％以下とする。 V: 0.005 to 0.5%
V is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when V is contained, the V content is preferably 0.005% or more. On the other hand, when the V content exceeds 0.5%, the generation of V-based precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when V is contained, the V content is preferably 0.5% or less. The V content is more preferably 0.01% or more. Moreover, the V content is more preferably 0.1% or less.
Ｖは、鋼板の焼き入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｖを含有する場合、Ｖ含有量を０．００５％以上とすることが好ましい。一方、Ｖの含有量が０．５％を超えると、Ｖ系析出物の生成が過剰となり、後加熱後の靭性低下を招く場合があるため、Ｖを含有する場合、Ｖ含有量を０．５％以下とすることが好ましい。Ｖ含有量は、より好ましくは０．０１％以上とする。また、Ｖ含有量は、より好ましくは０．１％以下とする。 V: 0.005 to 0.5%
V is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when V is contained, the V content is preferably 0.005% or more. On the other hand, when the V content exceeds 0.5%, the generation of V-based precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when V is contained, the V content is preferably 0.5% or less. The V content is more preferably 0.01% or more. Moreover, the V content is more preferably 0.1% or less.
Ｂ：０．０００２～０．００５０％
Ｂは、鋼板の焼入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｂを含有する場合、Ｂ含有量を０．０００２％以上とすることが好ましい。一方、Ｂ含有量が０．００５０％を超えるとＢ系化合物が増加して、後加熱後の靭性や耐遅れ破壊特性が低下する場合がある。したがって、Ｂを含有する場合、Ｂ含有量を０．００５０％以下とすることが好ましい。Ｂ含有量は、より好ましくは０．０００５％以上とする。また、Ｂ含有量は、より好ましくは０．００４０％以下とする。 B: 0.0002 to 0.0050%
B is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when B is contained, the B content is preferably 0.0002% or more. On the other hand, when the B content exceeds 0.0050%, B-based compounds increase, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when B is contained, the B content is preferably 0.0050% or less. The B content is more preferably 0.0005% or more. Moreover, the B content is more preferably 0.0040% or less.
Ｂは、鋼板の焼入れ性を高め、下部ベイナイトやマルテンサイトを生成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、Ｂを含有する場合、Ｂ含有量を０．０００２％以上とすることが好ましい。一方、Ｂ含有量が０．００５０％を超えるとＢ系化合物が増加して、後加熱後の靭性や耐遅れ破壊特性が低下する場合がある。したがって、Ｂを含有する場合、Ｂ含有量を０．００５０％以下とすることが好ましい。Ｂ含有量は、より好ましくは０．０００５％以上とする。また、Ｂ含有量は、より好ましくは０．００４０％以下とする。 B: 0.0002 to 0.0050%
B is an element effective in improving the hardenability of the steel sheet and generating lower bainite and martensite. In order to obtain such effects, when B is contained, the B content is preferably 0.0002% or more. On the other hand, when the B content exceeds 0.0050%, B-based compounds increase, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when B is contained, the B content is preferably 0.0050% or less. The B content is more preferably 0.0005% or more. Moreover, the B content is more preferably 0.0040% or less.
Ｃａ：０．０００１～０．００５０％、ＲＥＭ：０．０００１～０．００５０％
Ｃａ、ＲＥＭ（希土類元素）はそれぞれ、介在物の形態制御により後加熱後の靭性や耐遅れ破壊特性の向上に有効な元素である。このような効果を得るため、Ｃａ、ＲＥＭを含有する場合、それぞれの含有量を０．０００１％以上とすることが好ましい。一方、Ｃａ、ＲＥＭの含有量がそれぞれ０．００５０％を超えると、介在物量の増加の影響が過剰となり、後加熱後の靭性や耐遅れ破壊特性が低下する場合がある。よって、Ｃａ、ＲＥＭを含有する場合、Ｃａ、ＲＥＭの含有量はそれぞれ０．００５０％以下とすることが好ましい。Ｃａ含有量は、より好ましくは０．０００５％以上とする。また、Ｃａ含有量は、より好ましくは０．００３０％以下とする。ＲＥＭ含有量は、より好ましくは０．０００５％以上とする。また、ＲＥＭ含有量は、より好ましくは０．００３０％以下とする。なお、ＲＥＭは、Ｓｃ、Ｙと、原子番号５７のランタン（Ｌａ）から原子番号７１のルテチウム（Ｌｕ）までの１５元素の総称であり、ここでいうＲＥＭ含有量は、これらの元素の合計含有量である。 Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050%
Ca and REM (rare earth elements) are each an element effective in improving toughness and delayed fracture resistance after post-heating by controlling the shape of inclusions. In order to obtain such effects, when Ca and REM are contained, the respective contents are preferably 0.0001% or more. On the other hand, when the Ca and REM contents each exceed 0.0050%, the influence of the increase in the amount of inclusions becomes excessive, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when Ca and REM are contained, the Ca and REM contents are preferably 0.0050% or less. The Ca content is more preferably 0.0005% or more. Moreover, the Ca content is more preferably 0.0030% or less. The REM content is more preferably 0.0005% or more. Moreover, the REM content is more preferably 0.0030% or less. Note that REM is a collective term for Sc, Y, and 15 other elements ranging from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.
Ｃａ、ＲＥＭ（希土類元素）はそれぞれ、介在物の形態制御により後加熱後の靭性や耐遅れ破壊特性の向上に有効な元素である。このような効果を得るため、Ｃａ、ＲＥＭを含有する場合、それぞれの含有量を０．０００１％以上とすることが好ましい。一方、Ｃａ、ＲＥＭの含有量がそれぞれ０．００５０％を超えると、介在物量の増加の影響が過剰となり、後加熱後の靭性や耐遅れ破壊特性が低下する場合がある。よって、Ｃａ、ＲＥＭを含有する場合、Ｃａ、ＲＥＭの含有量はそれぞれ０．００５０％以下とすることが好ましい。Ｃａ含有量は、より好ましくは０．０００５％以上とする。また、Ｃａ含有量は、より好ましくは０．００３０％以下とする。ＲＥＭ含有量は、より好ましくは０．０００５％以上とする。また、ＲＥＭ含有量は、より好ましくは０．００３０％以下とする。なお、ＲＥＭは、Ｓｃ、Ｙと、原子番号５７のランタン（Ｌａ）から原子番号７１のルテチウム（Ｌｕ）までの１５元素の総称であり、ここでいうＲＥＭ含有量は、これらの元素の合計含有量である。 Ca: 0.0001 to 0.0050%, REM: 0.0001 to 0.0050%
Ca and REM (rare earth elements) are each an element effective in improving toughness and delayed fracture resistance after post-heating by controlling the shape of inclusions. In order to obtain such effects, when Ca and REM are contained, the respective contents are preferably 0.0001% or more. On the other hand, when the Ca and REM contents each exceed 0.0050%, the influence of the increase in the amount of inclusions becomes excessive, and the toughness and delayed fracture resistance after post-heating may decrease. Therefore, when Ca and REM are contained, the Ca and REM contents are preferably 0.0050% or less. The Ca content is more preferably 0.0005% or more. Moreover, the Ca content is more preferably 0.0030% or less. The REM content is more preferably 0.0005% or more. Moreover, the REM content is more preferably 0.0030% or less. Note that REM is a collective term for Sc, Y, and 15 other elements ranging from lanthanum (La) with atomic number 57 to lutetium (Lu) with atomic number 71, and the REM content referred to here is the total content of these elements.
Ｃｕ：０．００５～０．５％、Ｓｂ：０．００１０～０．１０％、Ｓｎ：０．００１０～０．１０％
Ｃｕ、Ｓｂ、Ｓｎはそれぞれ、腐食反応を遅らせ、後加熱後の耐遅れ破壊特性の向上に有効な元素である。このような効果を得るため、Ｃｕ、Ｓｂ、Ｓｎを含有する場合、それぞれ、Ｃｕ含有量を０．００５％以上、Ｓｂ含有量を０．００１０％以上、Ｓｎ含有量を０．００１０％以上とすることが好ましい。一方、Ｃｕの含有量が０．５％を超えると、Ｃｕ析出物の生成が過剰となり、後加熱後の靭性低下を招く場合があるため、Ｃｕを含有する場合、Ｃｕ含有量を０．５％以下とすることが好ましい。また、Ｓｂ、Ｓｎのそれぞれの含有量が０．１０％を超えると粒界脆化効果が過剰となって耐遅れ破壊特性が低下する場合があるため、Ｓｂ、Ｓｎを含有する場合、Ｓｂ、Ｓｎの含有量をそれぞれ０．１０％以下とすることが好ましい。Ｃｕ含有量は、より好ましくは０．０５％以上とする。また、Ｃｕ含有量は、より好ましくは０．３％以下とする。Ｓｂ含有量は、より好ましくは０．００５０％以上とする。また、Ｓｂ含有量は、より好ましくは０．０５０％以下とする。Ｓｎ含有量は、より好ましくは０．００５０％以上とする。また、Ｓｎ含有量は、より好ましくは０．０５０％以下とする。 Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, Sn: 0.0010 to 0.10%
Cu, Sb, and Sn are each an element that is effective in retarding the corrosion reaction and improving the delayed fracture resistance after post-heating. In order to obtain such effects, when Cu, Sb, and Sn are contained, it is preferable that the Cu content is 0.005% or more, the Sb content is 0.0010% or more, and the Sn content is 0.0010% or more, respectively. On the other hand, when the Cu content exceeds 0.5%, the generation of Cu precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when Cu is contained, it is preferable that the Cu content is 0.5% or less. Furthermore, when the Sb and Sn contents each exceed 0.10%, the grain boundary embrittlement effect becomes excessive, which may lead to a decrease in delayed fracture resistance, so when Sb and Sn are contained, it is preferable that the Sb and Sn contents each are 0.10% or less. The Cu content is more preferably 0.05% or more. Furthermore, the Cu content is more preferably 0.3% or less. The Sb content is more preferably 0.0050% or more. The Sb content is more preferably 0.050% or less. The Sn content is more preferably 0.0050% or more. The Sn content is more preferably 0.050% or less.
Ｃｕ、Ｓｂ、Ｓｎはそれぞれ、腐食反応を遅らせ、後加熱後の耐遅れ破壊特性の向上に有効な元素である。このような効果を得るため、Ｃｕ、Ｓｂ、Ｓｎを含有する場合、それぞれ、Ｃｕ含有量を０．００５％以上、Ｓｂ含有量を０．００１０％以上、Ｓｎ含有量を０．００１０％以上とすることが好ましい。一方、Ｃｕの含有量が０．５％を超えると、Ｃｕ析出物の生成が過剰となり、後加熱後の靭性低下を招く場合があるため、Ｃｕを含有する場合、Ｃｕ含有量を０．５％以下とすることが好ましい。また、Ｓｂ、Ｓｎのそれぞれの含有量が０．１０％を超えると粒界脆化効果が過剰となって耐遅れ破壊特性が低下する場合があるため、Ｓｂ、Ｓｎを含有する場合、Ｓｂ、Ｓｎの含有量をそれぞれ０．１０％以下とすることが好ましい。Ｃｕ含有量は、より好ましくは０．０５％以上とする。また、Ｃｕ含有量は、より好ましくは０．３％以下とする。Ｓｂ含有量は、より好ましくは０．００５０％以上とする。また、Ｓｂ含有量は、より好ましくは０．０５０％以下とする。Ｓｎ含有量は、より好ましくは０．００５０％以上とする。また、Ｓｎ含有量は、より好ましくは０．０５０％以下とする。 Cu: 0.005 to 0.5%, Sb: 0.0010 to 0.10%, Sn: 0.0010 to 0.10%
Cu, Sb, and Sn are each an element that is effective in retarding the corrosion reaction and improving the delayed fracture resistance after post-heating. In order to obtain such effects, when Cu, Sb, and Sn are contained, it is preferable that the Cu content is 0.005% or more, the Sb content is 0.0010% or more, and the Sn content is 0.0010% or more, respectively. On the other hand, when the Cu content exceeds 0.5%, the generation of Cu precipitates becomes excessive, which may lead to a decrease in toughness after post-heating, so when Cu is contained, it is preferable that the Cu content is 0.5% or less. Furthermore, when the Sb and Sn contents each exceed 0.10%, the grain boundary embrittlement effect becomes excessive, which may lead to a decrease in delayed fracture resistance, so when Sb and Sn are contained, it is preferable that the Sb and Sn contents each are 0.10% or less. The Cu content is more preferably 0.05% or more. Furthermore, the Cu content is more preferably 0.3% or less. The Sb content is more preferably 0.0050% or more. The Sb content is more preferably 0.050% or less. The Sn content is more preferably 0.0050% or more. The Sn content is more preferably 0.050% or less.
なお、Ｃｒ、Ｎｉ、Ｍｏ、Ｖ、Ｂ、Ｃａ、ＲＥＭ、Ｃｕ、Ｓｂ、Ｓｎの含有量が、上記の下限値未満であっても、本発明の効果を害さない。したがって、これらの成分の含有量が上記下限値未満の場合は、これらの元素を不可避的不純物として含むものとして扱う。また、本発明では、前記成分組成に加えて、さらに、質量％で、Ｍｇ、Ａｓ、Ｗ、Ｔａ、Ｐｂ、Ｚｒ、Ｈｆ、Ｔｅ、Ｂｉ、Ｓｅの１種または２種以上を合計で０．３％以下の範囲で含有させても良い。なお、これらの元素の含有量はそれぞれ０．０３％以下に制限することが好ましい。
The effect of the present invention is not impaired even if the content of Cr, Ni, Mo, V, B, Ca, REM, Cu, Sb, and Sn is less than the lower limit value. Therefore, when the content of these components is less than the lower limit value, these elements are treated as unavoidable impurities. In addition to the above-mentioned composition, the present invention may further contain one or more of Mg, As, W, Ta, Pb, Zr, Hf, Te, Bi, and Se in a total amount of 0.3% or less by mass. It is preferable to limit the content of each of these elements to 0.03% or less.
続いて、本発明の高強度熱延鋼板の鋼組織について説明する。
Next, we will explain the steel structure of the high-strength hot-rolled steel sheet of the present invention.
本発明の高強度熱延鋼板の鋼組織は、マルテンサイトおよび／または下部ベイナイトを主相とし、残留γが体積率で３％未満である。
The steel structure of the high-strength hot-rolled steel sheet of the present invention has martensite and/or lower bainite as the main phase, and the volume fraction of residual γ is less than 3%.
主相：マルテンサイトおよび／または下部ベイナイト
本発明では、高強度と後加熱後の優れた靭性および耐遅れ破壊特性を得るため、マルテンサイトおよび／または下部ベイナイトを主相とする組織とする。フェライトやパーライトや残留γ等が主相となると、高強度と後加熱後の優れた靭性や耐遅れ破壊特性の両立が困難となる。したがって、鋼組織はマルテンサイトおよび／または下部ベイナイトを主相とする。なお、マルテンサイトは自己焼戻し（オートテンパード）マルテンサイト、焼戻しマルテンサイトのいずれであっても構わないが、内部に炭化物を有さないフレッシュマルテンサイトは除く。また、下部ベイナイトは、焼き戻した下部ベイナイトであっても構わない。なお、本発明において、主相とは、面積率で５０％以上を占める相を意味する。主相の面積率は、６０％以上が好ましく、７５％以上がより好ましい。なお、本発明では、マルテンサイトが主相であってもよいし、下部ベイナイトが主相であってもよいし、マルテンサイトと下部ベイナイトの合計が主相であってもよい。主相の面積率の上限は、特に限定されず、１００％であってもよい。主相の面積率は、一例としては、１００％未満であってもよく、９８％以下であってもよい。 Main phase: martensite and/or lower bainite In the present invention, in order to obtain high strength and excellent toughness and delayed fracture resistance after post-heating, the microstructure is made to have martensite and/or lower bainite as the main phase. If ferrite, pearlite, residual γ, etc. become the main phase, it becomes difficult to achieve both high strength and excellent toughness and delayed fracture resistance after post-heating. Therefore, the steel microstructure is made to have martensite and/or lower bainite as the main phase. Note that the martensite may be either auto-tempered martensite or tempered martensite, but fresh martensite having no carbides inside is excluded. Also, the lower bainite may be tempered lower bainite. Note that in the present invention, the main phase means a phase that occupies 50% or more in terms of area ratio. The area ratio of the main phase is preferably 60% or more, and more preferably 75% or more. Note that in the present invention, martensite may be the main phase, lower bainite may be the main phase, or the total of martensite and lower bainite may be the main phase. The upper limit of the area ratio of the main phase is not particularly limited and may be 100%. The area ratio of the main phase may be, for example, less than 100% or 98% or less.
本発明では、高強度と後加熱後の優れた靭性および耐遅れ破壊特性を得るため、マルテンサイトおよび／または下部ベイナイトを主相とする組織とする。フェライトやパーライトや残留γ等が主相となると、高強度と後加熱後の優れた靭性や耐遅れ破壊特性の両立が困難となる。したがって、鋼組織はマルテンサイトおよび／または下部ベイナイトを主相とする。なお、マルテンサイトは自己焼戻し（オートテンパード）マルテンサイト、焼戻しマルテンサイトのいずれであっても構わないが、内部に炭化物を有さないフレッシュマルテンサイトは除く。また、下部ベイナイトは、焼き戻した下部ベイナイトであっても構わない。なお、本発明において、主相とは、面積率で５０％以上を占める相を意味する。主相の面積率は、６０％以上が好ましく、７５％以上がより好ましい。なお、本発明では、マルテンサイトが主相であってもよいし、下部ベイナイトが主相であってもよいし、マルテンサイトと下部ベイナイトの合計が主相であってもよい。主相の面積率の上限は、特に限定されず、１００％であってもよい。主相の面積率は、一例としては、１００％未満であってもよく、９８％以下であってもよい。 Main phase: martensite and/or lower bainite In the present invention, in order to obtain high strength and excellent toughness and delayed fracture resistance after post-heating, the microstructure is made to have martensite and/or lower bainite as the main phase. If ferrite, pearlite, residual γ, etc. become the main phase, it becomes difficult to achieve both high strength and excellent toughness and delayed fracture resistance after post-heating. Therefore, the steel microstructure is made to have martensite and/or lower bainite as the main phase. Note that the martensite may be either auto-tempered martensite or tempered martensite, but fresh martensite having no carbides inside is excluded. Also, the lower bainite may be tempered lower bainite. Note that in the present invention, the main phase means a phase that occupies 50% or more in terms of area ratio. The area ratio of the main phase is preferably 60% or more, and more preferably 75% or more. Note that in the present invention, martensite may be the main phase, lower bainite may be the main phase, or the total of martensite and lower bainite may be the main phase. The upper limit of the area ratio of the main phase is not particularly limited and may be 100%. The area ratio of the main phase may be, for example, less than 100% or 98% or less.
残留オーステナイト（残留γ）量：３％未満
残留オーステナイト（残留γ）は、後加熱後はパーライトに変態することで強度と靭性を著しく低下させる組織であるため極力低減することが好ましい。本発明では、体積率で、残留γが３％未満まで許容される。よって、残留γは体積率で３％未満とする。残留γは、体積率で、好ましくは２％未満であり、より好ましくは１％未満である。残留γの体積率の下限は特に限定されず、残留γの体積率は０％であってもよい。 Amount of retained austenite (residual γ): less than 3% Since retained austenite (residual γ) is a structure that significantly reduces strength and toughness by transforming into pearlite after post-heating, it is preferable to reduce it as much as possible. In the present invention, the volume fraction of retained γ is allowed to be less than 3%. Therefore, the volume fraction of retained γ is set to less than 3%. The volume fraction of retained γ is preferably less than 2%, and more preferably less than 1%. There is no particular limit on the lower limit of the volume fraction of retained γ, and the volume fraction of retained γ may be 0%.
残留オーステナイト（残留γ）は、後加熱後はパーライトに変態することで強度と靭性を著しく低下させる組織であるため極力低減することが好ましい。本発明では、体積率で、残留γが３％未満まで許容される。よって、残留γは体積率で３％未満とする。残留γは、体積率で、好ましくは２％未満であり、より好ましくは１％未満である。残留γの体積率の下限は特に限定されず、残留γの体積率は０％であってもよい。 Amount of retained austenite (residual γ): less than 3% Since retained austenite (residual γ) is a structure that significantly reduces strength and toughness by transforming into pearlite after post-heating, it is preferable to reduce it as much as possible. In the present invention, the volume fraction of retained γ is allowed to be less than 3%. Therefore, the volume fraction of retained γ is set to less than 3%. The volume fraction of retained γ is preferably less than 2%, and more preferably less than 1%. There is no particular limit on the lower limit of the volume fraction of retained γ, and the volume fraction of retained γ may be 0%.
なお、マルテンサイト、下部ベイナイト、残留γ以外の相（その他の相）としては、フェライト、パーライト、上部ベイナイトのうちの１種または２種以上が挙げられる。その他の相の面積率は、合計で３０％以下が好ましく、合計で２５％以下がより好ましい。その他の相の面積率の下限は、特に限定されず、その他の相の面積率は、合計で０％であってもよい。
The phases other than martensite, lower bainite, and residual gamma (other phases) may be one or more of ferrite, pearlite, and upper bainite. The total area ratio of the other phases is preferably 30% or less, and more preferably 25% or less. There is no particular lower limit to the area ratio of the other phases, and the total area ratio of the other phases may be 0%.
（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）：０．３００以上０．８００未満
Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である［（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）］が０．３００未満では、後加熱時に析出物となって強度低下を相殺するための固溶Ｔｉ量および固溶Ｎｂ量が不十分となる。その結果、後加熱後の優れた強度が得られなくなる。一方、０．８００以上では、後加熱後の析出物析出による強度上昇が過剰となり、後加熱後の優れた靭性が得られなくなる。したがって、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）は０．３００以上０．８００未満とする。好ましくは０．３５０以上とする。また、好ましくは０．７００以下とする。なお、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）は、実施例に記載の方法により求められる。 (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount): 0.300 or more and less than 0.800 If the ratio of the total of the dissolved Ti amount and the dissolved Nb amount to the total of the Ti content and the Nb content, [(Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount)], is less than 0.300, the amount of dissolved Ti and the amount of dissolved Nb that become precipitates during post-heating to offset the strength reduction will be insufficient. As a result, excellent strength after post-heating will not be obtained. On the other hand, if it is 0.800 or more, the strength increase due to the precipitation of precipitates after post-heating will be excessive, and excellent toughness after post-heating will not be obtained. Therefore, (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) is 0.300 or more and less than 0.800. It is preferably 0.350 or more. It is also preferably 0.700 or less. The value of (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb) is determined by the method described in the Examples.
Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である［（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）］が０．３００未満では、後加熱時に析出物となって強度低下を相殺するための固溶Ｔｉ量および固溶Ｎｂ量が不十分となる。その結果、後加熱後の優れた強度が得られなくなる。一方、０．８００以上では、後加熱後の析出物析出による強度上昇が過剰となり、後加熱後の優れた靭性が得られなくなる。したがって、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）は０．３００以上０．８００未満とする。好ましくは０．３５０以上とする。また、好ましくは０．７００以下とする。なお、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）は、実施例に記載の方法により求められる。 (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount): 0.300 or more and less than 0.800 If the ratio of the total of the dissolved Ti amount and the dissolved Nb amount to the total of the Ti content and the Nb content, [(Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount)], is less than 0.300, the amount of dissolved Ti and the amount of dissolved Nb that become precipitates during post-heating to offset the strength reduction will be insufficient. As a result, excellent strength after post-heating will not be obtained. On the other hand, if it is 0.800 or more, the strength increase due to the precipitation of precipitates after post-heating will be excessive, and excellent toughness after post-heating will not be obtained. Therefore, (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) is 0.300 or more and less than 0.800. It is preferably 0.350 or more. It is also preferably 0.700 or less. The value of (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb) is determined by the method described in the Examples.
粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計：０．０１０～０．０３０質量％
粒径１００ｎｍ以上のＴｉ含有析出物およびＮｂ含有析出物を一定以上含有させることで、後加熱の際に該析出物の成長と新たなＴｉＣおよびＮｂＣ等の析出が競合する。これにより、微細なＴｉＣやＮｂＣ等の析出が適度に抑制され、過度な強度上昇や靭性の低下を抑制することができる。このような効果を得るには粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０質量％以上とする必要がある。一方、前記Ｔｉ量とＮｂ量の合計が０．０３０質量％を超えると粗大析出物による靭性の低下が顕著になるため、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０３０質量％以下とする必要がある。したがって、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０～０．０３０質量％とする。好ましくは、０．０１３質量％以上とする。また、好ましくは０．０２７質量％以下とする。なお、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計は、実施例に記載の方法により求められる。 The total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more: 0.010 to 0.030 mass%
By containing a certain amount of Ti-containing precipitates and Nb-containing precipitates having a grain size of 100 nm or more, the growth of the precipitates competes with the precipitation of new TiC, NbC, etc. during post-heating. This appropriately suppresses the precipitation of fine TiC, NbC, etc., and can suppress excessive strength increase and toughness decrease. To obtain such an effect, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 mass% or more. On the other hand, if the total amount of Ti and Nb exceeds 0.030 mass%, the decrease in toughness due to coarse precipitates becomes significant, so the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.030 mass% or less. Therefore, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 to 0.030 mass%. Preferably, it is 0.013 mass% or more. Also, it is preferably 0.027 mass% or less. The total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is determined by the method described in the Examples.
粒径１００ｎｍ以上のＴｉ含有析出物およびＮｂ含有析出物を一定以上含有させることで、後加熱の際に該析出物の成長と新たなＴｉＣおよびＮｂＣ等の析出が競合する。これにより、微細なＴｉＣやＮｂＣ等の析出が適度に抑制され、過度な強度上昇や靭性の低下を抑制することができる。このような効果を得るには粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０質量％以上とする必要がある。一方、前記Ｔｉ量とＮｂ量の合計が０．０３０質量％を超えると粗大析出物による靭性の低下が顕著になるため、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０３０質量％以下とする必要がある。したがって、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０～０．０３０質量％とする。好ましくは、０．０１３質量％以上とする。また、好ましくは０．０２７質量％以下とする。なお、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計は、実施例に記載の方法により求められる。 The total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more: 0.010 to 0.030 mass%
By containing a certain amount of Ti-containing precipitates and Nb-containing precipitates having a grain size of 100 nm or more, the growth of the precipitates competes with the precipitation of new TiC, NbC, etc. during post-heating. This appropriately suppresses the precipitation of fine TiC, NbC, etc., and can suppress excessive strength increase and toughness decrease. To obtain such an effect, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 mass% or more. On the other hand, if the total amount of Ti and Nb exceeds 0.030 mass%, the decrease in toughness due to coarse precipitates becomes significant, so the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.030 mass% or less. Therefore, the total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more must be 0.010 to 0.030 mass%. Preferably, it is 0.013 mass% or more. Also, it is preferably 0.027 mass% or less. The total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is determined by the method described in the Examples.
表面から板厚中央方向に１００μｍまでの表層領域における｛１１０｝＜１１１＞方位の極密度：１．８～５．０
鋼板の表面から板厚中央方向に１００μｍまでの表層領域は、打抜きや高速変形時の破面形成に強く影響する。この領域の｛１１０｝＜１１１＞方位の極密度を１．８～５．０の範囲に制御することで、後加熱後に優れた靭性を得ることができ、また、打抜きの破面性状が良好となり、後加熱後に優れた耐遅れ破壊特性を得ることができる。このような効果を得るには、表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度を１．８以上とする必要がある。一方、前記極密度が５．０を超えると後加熱後の強度低下が顕著となり、後加熱後の優れた強度（５０以下のΔＨＶ）が得られなくなる。したがって、表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度を１．８～５．０とする。好ましくは２．０以上とする。また、好ましくは４．０以下とし、より好ましくは３．０以下とする。なお、表面から板厚中央方向に１００μｍまでの表層領域における｛１１０｝＜１１１＞方位の極密度は、実施例に記載の方法により求められる。 Pole density of {110}<111> orientation in the surface layer region from the surface to 100 μm in the center direction of the sheet thickness: 1.8 to 5.0
The surface layer region from the surface of the steel plate to 100 μm in the thickness center direction strongly influences the formation of fracture surface during punching or high-speed deformation. By controlling the pole density of the {110}<111> orientation in this region to the range of 1.8 to 5.0, excellent toughness can be obtained after post-heating, and the fracture surface properties of punching become good, and excellent delayed fracture resistance properties can be obtained after post-heating. To obtain such effects, the pole density of the {110}<111> orientation needs to be 1.8 or more in the surface layer region from the surface to 100 μm in the thickness center direction. On the other hand, if the pole density exceeds 5.0, the strength reduction after post-heating becomes significant, and excellent strength after post-heating (ΔHV of 50 or less) cannot be obtained. Therefore, the pole density of the {110}<111> orientation is set to 1.8 to 5.0 in the surface layer region from the surface to 100 μm in the thickness center direction. It is preferably set to 2.0 or more. It is also preferably set to 4.0 or less, and more preferably set to 3.0 or less. The pole density of the {110}<111> orientation in the surface layer region extending from the surface to 100 μm in the sheet thickness center direction is determined by the method described in the examples.
鋼板の表面から板厚中央方向に１００μｍまでの表層領域は、打抜きや高速変形時の破面形成に強く影響する。この領域の｛１１０｝＜１１１＞方位の極密度を１．８～５．０の範囲に制御することで、後加熱後に優れた靭性を得ることができ、また、打抜きの破面性状が良好となり、後加熱後に優れた耐遅れ破壊特性を得ることができる。このような効果を得るには、表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度を１．８以上とする必要がある。一方、前記極密度が５．０を超えると後加熱後の強度低下が顕著となり、後加熱後の優れた強度（５０以下のΔＨＶ）が得られなくなる。したがって、表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度を１．８～５．０とする。好ましくは２．０以上とする。また、好ましくは４．０以下とし、より好ましくは３．０以下とする。なお、表面から板厚中央方向に１００μｍまでの表層領域における｛１１０｝＜１１１＞方位の極密度は、実施例に記載の方法により求められる。 Pole density of {110}<111> orientation in the surface layer region from the surface to 100 μm in the center direction of the sheet thickness: 1.8 to 5.0
The surface layer region from the surface of the steel plate to 100 μm in the thickness center direction strongly influences the formation of fracture surface during punching or high-speed deformation. By controlling the pole density of the {110}<111> orientation in this region to the range of 1.8 to 5.0, excellent toughness can be obtained after post-heating, and the fracture surface properties of punching become good, and excellent delayed fracture resistance properties can be obtained after post-heating. To obtain such effects, the pole density of the {110}<111> orientation needs to be 1.8 or more in the surface layer region from the surface to 100 μm in the thickness center direction. On the other hand, if the pole density exceeds 5.0, the strength reduction after post-heating becomes significant, and excellent strength after post-heating (ΔHV of 50 or less) cannot be obtained. Therefore, the pole density of the {110}<111> orientation is set to 1.8 to 5.0 in the surface layer region from the surface to 100 μm in the thickness center direction. It is preferably set to 2.0 or more. It is also preferably set to 4.0 or less, and more preferably set to 3.0 or less. The pole density of the {110}<111> orientation in the surface layer region extending from the surface to 100 μm in the sheet thickness center direction is determined by the method described in the examples.
＜高強度熱延鋼板の製造方法＞
本発明の高強度熱延鋼板は、上記成分組成を有するスラブを１１５０～１３００℃の温度域に加熱し、該温度域で０．２～３．５時間保持し、次いで、熱間圧延を施すに際し、１０８０℃以上の温度域での合計圧下率が８０～９０％、９００℃以下の温度域での合計圧下率が２０％以上、かつ、下記式で求められるＴ（℃）以下での１パスあたりの圧下率が２５％以下となる条件で圧延した後、１．０ｓ以上放冷し、次いで、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却し、５５０℃に達してから急冷を開始するまでの時間を０．５～４．０ｓとし、次いで、１００～２５０℃の巻取り温度までを２００℃／ｓ以上の冷却速度で急冷し、前記巻取り温度で巻取ることにより製造する。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。
また、１０８０℃以上の温度域での合計圧下率は、熱間圧延を施す前のスラブの厚さを基準とし、これと１０８０℃時点の板厚との比から求める。また、９００℃以下の温度域での合計圧下率は、９００℃時点の板厚を基準とし、これと最終板厚との比から求める。また、Ｔ（℃）以下での１パスあたりの圧下率は、Ｔ（℃）以下での各パスの圧延前後の板厚の比から求める。 <Method of manufacturing high-strength hot-rolled steel sheet>
The high-strength hot-rolled steel sheet of the present invention is produced by heating a slab having the above-mentioned composition to a temperature range of 1150 to 1300°C, holding the slab in the temperature range for 0.2 to 3.5 hours, and then hot rolling the slab under conditions in which the total reduction in the temperature range of 1080°C or higher is 80 to 90%, the total reduction in the temperature range of 900°C or lower is 20% or more, and the reduction per pass at or below T (°C) calculated by the following formula is 25% or less, followed by allowing the slab to cool for 1.0 s or more, then cooling the slab at a temperature range up to 550°C at an average cooling rate of 50°C/s or more, setting the time from reaching 550°C to the start of quenching to 0.5 to 4.0 s, then quenching the slab to a coiling temperature of 100 to 250°C at a cooling rate of 200°C/s or more, and coiling the slab at the coiling temperature.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
The total reduction in the temperature range of 1080° C. or higher is determined from the ratio of the slab thickness before hot rolling to the plate thickness at 1080° C. The total reduction in the temperature range of 900° C. or lower is determined from the ratio of the plate thickness at 900° C. to the final plate thickness. The reduction per pass at or below T (° C.) is determined from the ratio of the plate thickness before and after each pass of rolling at or below T (° C.).
本発明の高強度熱延鋼板は、上記成分組成を有するスラブを１１５０～１３００℃の温度域に加熱し、該温度域で０．２～３．５時間保持し、次いで、熱間圧延を施すに際し、１０８０℃以上の温度域での合計圧下率が８０～９０％、９００℃以下の温度域での合計圧下率が２０％以上、かつ、下記式で求められるＴ（℃）以下での１パスあたりの圧下率が２５％以下となる条件で圧延した後、１．０ｓ以上放冷し、次いで、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却し、５５０℃に達してから急冷を開始するまでの時間を０．５～４．０ｓとし、次いで、１００～２５０℃の巻取り温度までを２００℃／ｓ以上の冷却速度で急冷し、前記巻取り温度で巻取ることにより製造する。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。
また、１０８０℃以上の温度域での合計圧下率は、熱間圧延を施す前のスラブの厚さを基準とし、これと１０８０℃時点の板厚との比から求める。また、９００℃以下の温度域での合計圧下率は、９００℃時点の板厚を基準とし、これと最終板厚との比から求める。また、Ｔ（℃）以下での１パスあたりの圧下率は、Ｔ（℃）以下での各パスの圧延前後の板厚の比から求める。 <Method of manufacturing high-strength hot-rolled steel sheet>
The high-strength hot-rolled steel sheet of the present invention is produced by heating a slab having the above-mentioned composition to a temperature range of 1150 to 1300°C, holding the slab in the temperature range for 0.2 to 3.5 hours, and then hot rolling the slab under conditions in which the total reduction in the temperature range of 1080°C or higher is 80 to 90%, the total reduction in the temperature range of 900°C or lower is 20% or more, and the reduction per pass at or below T (°C) calculated by the following formula is 25% or less, followed by allowing the slab to cool for 1.0 s or more, then cooling the slab at a temperature range up to 550°C at an average cooling rate of 50°C/s or more, setting the time from reaching 550°C to the start of quenching to 0.5 to 4.0 s, then quenching the slab to a coiling temperature of 100 to 250°C at a cooling rate of 200°C/s or more, and coiling the slab at the coiling temperature.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
The total reduction in the temperature range of 1080° C. or higher is determined from the ratio of the slab thickness before hot rolling to the plate thickness at 1080° C. The total reduction in the temperature range of 900° C. or lower is determined from the ratio of the plate thickness at 900° C. to the final plate thickness. The reduction per pass at or below T (° C.) is determined from the ratio of the plate thickness before and after each pass of rolling at or below T (° C.).
以下、詳しく説明する。なお、上記した温度は鋼板の幅中央部の表面の温度であり、上記した平均冷却速度、冷却速度は、それぞれ鋼板の幅中央部の表面の平均冷却速度、冷却速度である。また、平均冷却速度は、特に断らない限り、［（冷却開始温度－冷却停止温度）／冷却開始温度から冷却停止温度までの冷却時間］とする。
A detailed explanation is provided below. The above temperatures are the surface temperatures at the center of the width of the steel plate, and the above average cooling rates and cooling speeds are the average cooling rate and cooling speed at the surface at the center of the width of the steel plate, respectively. Furthermore, unless otherwise specified, the average cooling rate is [(cooling start temperature - cooling stop temperature) / cooling time from cooling start temperature to cooling stop temperature].
スラブの加熱温度：１１５０～１３００℃
スラブの加熱温度が１１５０℃未満では、Ｔｉ含有析出物の溶解が不十分となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。一方、スラブの加熱温度が１３００℃を超えると、Ｔｉ含有析出物やＮｂ含有析出物の溶解が過剰となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。したがって、スラブの加熱温度は１１５０～１３００℃とする。前記加熱温度は、好ましくは１１７０℃以上とし、より好ましくは１１８５℃以上とする。また、前記加熱温度は、好ましくは１２８０℃以下とし、より好ましくは１２６５℃以下とする。 Slab heating temperature: 1150-1300°C
If the heating temperature of the slab is less than 1150°C, the dissolution of Ti-containing precipitates is insufficient. As a result, the value of (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) of 0.300 or more and less than 0.800, or the value of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more of 0.010 to 0.030 mass% cannot be obtained. On the other hand, if the heating temperature of the slab exceeds 1300°C, the dissolution of Ti-containing precipitates and Nb-containing precipitates becomes excessive. As a result, the value of (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) of 0.300 or more and less than 0.800, or the value of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more of 0.010 to 0.030 mass% cannot be obtained. Therefore, the heating temperature of the slab is set to 1150 to 1300°C. The heating temperature is preferably 1170° C. or higher, and more preferably 1185° C. or higher. The heating temperature is preferably 1280° C. or lower, and more preferably 1265° C. or lower.
スラブの加熱温度が１１５０℃未満では、Ｔｉ含有析出物の溶解が不十分となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。一方、スラブの加熱温度が１３００℃を超えると、Ｔｉ含有析出物やＮｂ含有析出物の溶解が過剰となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。したがって、スラブの加熱温度は１１５０～１３００℃とする。前記加熱温度は、好ましくは１１７０℃以上とし、より好ましくは１１８５℃以上とする。また、前記加熱温度は、好ましくは１２８０℃以下とし、より好ましくは１２６５℃以下とする。 Slab heating temperature: 1150-1300°C
If the heating temperature of the slab is less than 1150°C, the dissolution of Ti-containing precipitates is insufficient. As a result, the value of (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) of 0.300 or more and less than 0.800, or the value of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more of 0.010 to 0.030 mass% cannot be obtained. On the other hand, if the heating temperature of the slab exceeds 1300°C, the dissolution of Ti-containing precipitates and Nb-containing precipitates becomes excessive. As a result, the value of (Solute Ti amount + Solute Nb amount) / (Total Ti amount + Total Nb amount) of 0.300 or more and less than 0.800, or the value of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more of 0.010 to 0.030 mass% cannot be obtained. Therefore, the heating temperature of the slab is set to 1150 to 1300°C. The heating temperature is preferably 1170° C. or higher, and more preferably 1185° C. or higher. The heating temperature is preferably 1280° C. or lower, and more preferably 1265° C. or lower.
１１５０～１３００℃の温度域での保持時間：０．２～３．５時間
１１５０～１３００℃の温度域での保持時間が０．２時間未満では、Ｔｉ含有析出物およびＮｂ含有析出物の溶解が不十分となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。一方、前記温度域での保持時間が３．５時間を超えると、表層近傍での脱炭が顕著になり、表層にフェライト、上部ベイナイト、残留γ等が生じやすくなり、本発明の組織が得られなくなる。したがって、スラブの前記温度域での保持時間は０．２～３．５時間とする。前記保持時間は、好ましくは０．４時間以上とする。また、前記保持時間は、好ましくは２．５時間以下とする。 Holding time in the temperature range of 1150 to 1300°C: 0.2 to 3.5 hours If the holding time in the temperature range of 1150 to 1300°C is less than 0.2 hours, the dissolution of Ti-containing precipitates and Nb-containing precipitates will be insufficient. As a result, a value of (amount of dissolved Ti + amount of dissolved Nb)/(total amount of Ti + total amount of Nb) of 0.300 or more and less than 0.800, or a value of 0.010 to 0.030 mass% of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more will not be obtained. On the other hand, if the holding time in the above temperature range exceeds 3.5 hours, decarburization in the vicinity of the surface layer will become significant, and ferrite, upper bainite, residual γ, etc. will be easily generated in the surface layer, and the structure of the present invention will not be obtained. Therefore, the holding time of the slab in the above temperature range is set to 0.2 to 3.5 hours. The holding time is preferably 0.4 hours or more. The holding time is preferably 2.5 hours or less.
１１５０～１３００℃の温度域での保持時間が０．２時間未満では、Ｔｉ含有析出物およびＮｂ含有析出物の溶解が不十分となる。その結果、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値や、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計の０．０１０～０．０３０質量％の値が得られなくなる。一方、前記温度域での保持時間が３．５時間を超えると、表層近傍での脱炭が顕著になり、表層にフェライト、上部ベイナイト、残留γ等が生じやすくなり、本発明の組織が得られなくなる。したがって、スラブの前記温度域での保持時間は０．２～３．５時間とする。前記保持時間は、好ましくは０．４時間以上とする。また、前記保持時間は、好ましくは２．５時間以下とする。 Holding time in the temperature range of 1150 to 1300°C: 0.2 to 3.5 hours If the holding time in the temperature range of 1150 to 1300°C is less than 0.2 hours, the dissolution of Ti-containing precipitates and Nb-containing precipitates will be insufficient. As a result, a value of (amount of dissolved Ti + amount of dissolved Nb)/(total amount of Ti + total amount of Nb) of 0.300 or more and less than 0.800, or a value of 0.010 to 0.030 mass% of the total amount of Ti and Nb present as precipitates with a grain size of 100 nm or more will not be obtained. On the other hand, if the holding time in the above temperature range exceeds 3.5 hours, decarburization in the vicinity of the surface layer will become significant, and ferrite, upper bainite, residual γ, etc. will be easily generated in the surface layer, and the structure of the present invention will not be obtained. Therefore, the holding time of the slab in the above temperature range is set to 0.2 to 3.5 hours. The holding time is preferably 0.4 hours or more. The holding time is preferably 2.5 hours or less.
１０８０℃以上の温度域での合計圧下率：８０～９０％
１０８０℃以上の温度域で合計圧下率８０～９０％の圧下を施すことで、１００ｎｍ以上の粗大Ｔｉ含有析出物やＮｂ含有析出物の生成と成長を促進することができる。その結果、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０～０．０３０質量％とすることができる。該合計圧下率が８０％未満では粒径１００ｎｍ以上の析出物の生成が不十分となり、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計が０．０１０質量％未満となる。一方、該合計圧下率が９０％を超えると、粒径１００ｎｍ以上の析出物の生成が過剰となり、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計が０．０３０質量％超となる。したがって、１０８０℃以上の温度域での合計圧下率は８０～９０％とする。前記合計圧下率は、好ましくは８１％以上とする。また、前記合計圧下率は、好ましくは８８％以下とする。 Total reduction rate at temperatures above 1080°C: 80-90%
By carrying out a total reduction of 80 to 90% in a temperature range of 1080°C or more, it is possible to promote the generation and growth of coarse Ti-containing precipitates and Nb-containing precipitates having a particle size of 100 nm or more. As a result, the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more can be set to 0.010 to 0.030 mass%. If the total reduction is less than 80%, the generation of precipitates having a particle size of 100 nm or more is insufficient, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more is less than 0.010 mass%. On the other hand, if the total reduction exceeds 90%, the generation of precipitates having a particle size of 100 nm or more is excessive, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more exceeds 0.030 mass%. Therefore, the total reduction in a temperature range of 1080°C or more is set to 80 to 90%. The total reduction is preferably set to 81% or more. The total rolling reduction is preferably 88% or less.
１０８０℃以上の温度域で合計圧下率８０～９０％の圧下を施すことで、１００ｎｍ以上の粗大Ｔｉ含有析出物やＮｂ含有析出物の生成と成長を促進することができる。その結果、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計を０．０１０～０．０３０質量％とすることができる。該合計圧下率が８０％未満では粒径１００ｎｍ以上の析出物の生成が不十分となり、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計が０．０１０質量％未満となる。一方、該合計圧下率が９０％を超えると、粒径１００ｎｍ以上の析出物の生成が過剰となり、粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量の合計が０．０３０質量％超となる。したがって、１０８０℃以上の温度域での合計圧下率は８０～９０％とする。前記合計圧下率は、好ましくは８１％以上とする。また、前記合計圧下率は、好ましくは８８％以下とする。 Total reduction rate at temperatures above 1080°C: 80-90%
By carrying out a total reduction of 80 to 90% in a temperature range of 1080°C or more, it is possible to promote the generation and growth of coarse Ti-containing precipitates and Nb-containing precipitates having a particle size of 100 nm or more. As a result, the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more can be set to 0.010 to 0.030 mass%. If the total reduction is less than 80%, the generation of precipitates having a particle size of 100 nm or more is insufficient, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more is less than 0.010 mass%. On the other hand, if the total reduction exceeds 90%, the generation of precipitates having a particle size of 100 nm or more is excessive, and the total amount of Ti and Nb present as precipitates having a particle size of 100 nm or more exceeds 0.030 mass%. Therefore, the total reduction in a temperature range of 1080°C or more is set to 80 to 90%. The total reduction is preferably set to 81% or more. The total rolling reduction is preferably 88% or less.
９００℃以下の温度域での合計圧下率が２０％以上
９００℃以下の合計圧下率が２０％未満となると、ひずみ誘起析出が抑制されてＴｉ含有析出物やＮｂ含有析出物が減少し、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値が得られなくなる。もしくは表層部の集合組織の発達が不十分となり、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、９００℃以下の温度域での合計圧下率は２０％以上とする。なお、前記合計圧下率の上限は、特に限定されないが、前記合計圧下率は、８０％以下が好ましく、６０％以下がより好ましい。 Total rolling reduction in the temperature range of 900°C or less is 20% or more If the total rolling reduction in the temperature range of 900°C or less is less than 20%, strain-induced precipitation is suppressed, Ti-containing precipitates and Nb-containing precipitates are reduced, and a value of (solubilized Ti amount + solid-solubilized Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. Or, the texture of the surface layer portion is insufficiently developed, and the pole density of the {110}<111> orientation in the surface layer region is not able to be 1.8 to 5.0. Therefore, the total rolling reduction in the temperature range of 900°C or less is set to 20% or more. The upper limit of the total rolling reduction is not particularly limited, but the total rolling reduction is preferably 80% or less, more preferably 60% or less.
９００℃以下の合計圧下率が２０％未満となると、ひずみ誘起析出が抑制されてＴｉ含有析出物やＮｂ含有析出物が減少し、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値が得られなくなる。もしくは表層部の集合組織の発達が不十分となり、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、９００℃以下の温度域での合計圧下率は２０％以上とする。なお、前記合計圧下率の上限は、特に限定されないが、前記合計圧下率は、８０％以下が好ましく、６０％以下がより好ましい。 Total rolling reduction in the temperature range of 900°C or less is 20% or more If the total rolling reduction in the temperature range of 900°C or less is less than 20%, strain-induced precipitation is suppressed, Ti-containing precipitates and Nb-containing precipitates are reduced, and a value of (solubilized Ti amount + solid-solubilized Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. Or, the texture of the surface layer portion is insufficiently developed, and the pole density of the {110}<111> orientation in the surface layer region is not able to be 1.8 to 5.0. Therefore, the total rolling reduction in the temperature range of 900°C or less is set to 20% or more. The upper limit of the total rolling reduction is not particularly limited, but the total rolling reduction is preferably 80% or less, more preferably 60% or less.
Ｔ（℃）以下での１パスあたりの圧下率：２５％以下
下記式で求められるＴ（℃）以下で１パスあたり２５％超えの圧下を施すと、ひずみ誘起析出が促進してＴｉ含有析出物やＮｂ含有析出物が増大し、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値が得られなくなる。同時に、表層部の集合組織が発達して、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、前記Ｔ（℃）以下での１パスあたりの圧下率が２５％以下とする。前記圧下率は、好ましくは２０％以下とし、より好ましくは１８％以下とする。前記圧下率の下限は特に限定されないが、５％以下では粗大粒が生じる場合があるため、前記圧下率は５％超とすることが好ましい。前記圧下率は、より好ましくは７％以上とする。
なお、Ｔ（℃）は下記式で求められる。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。 Reduction rate per pass at T (°C): 25% or less When reduction of more than 25% per pass is applied at T (°C) or less calculated by the following formula, strain-induced precipitation is promoted, Ti-containing precipitates and Nb-containing precipitates increase, and a value of (solute Ti amount + solute Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. At the same time, the texture of the surface layer part develops, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. Therefore, the reduction rate per pass at T (°C) or less is set to 25% or less. The reduction rate is preferably set to 20% or less, more preferably set to 18% or less. The lower limit of the reduction rate is not particularly limited, but since coarse grains may occur at 5% or less, the reduction rate is preferably set to more than 5%. The reduction rate is more preferably set to 7% or more.
Incidentally, T (°C) can be calculated by the following formula.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
下記式で求められるＴ（℃）以下で１パスあたり２５％超えの圧下を施すと、ひずみ誘起析出が促進してＴｉ含有析出物やＮｂ含有析出物が増大し、（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）の０．３００以上０．８００未満の値が得られなくなる。同時に、表層部の集合組織が発達して、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、前記Ｔ（℃）以下での１パスあたりの圧下率が２５％以下とする。前記圧下率は、好ましくは２０％以下とし、より好ましくは１８％以下とする。前記圧下率の下限は特に限定されないが、５％以下では粗大粒が生じる場合があるため、前記圧下率は５％超とすることが好ましい。前記圧下率は、より好ましくは７％以上とする。
なお、Ｔ（℃）は下記式で求められる。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。 Reduction rate per pass at T (°C): 25% or less When reduction of more than 25% per pass is applied at T (°C) or less calculated by the following formula, strain-induced precipitation is promoted, Ti-containing precipitates and Nb-containing precipitates increase, and a value of (solute Ti amount + solute Nb amount) / (total Ti amount + total Nb amount) of 0.300 or more and less than 0.800 cannot be obtained. At the same time, the texture of the surface layer part develops, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. Therefore, the reduction rate per pass at T (°C) or less is set to 25% or less. The reduction rate is preferably set to 20% or less, more preferably set to 18% or less. The lower limit of the reduction rate is not particularly limited, but since coarse grains may occur at 5% or less, the reduction rate is preferably set to more than 5%. The reduction rate is more preferably set to 7% or more.
Incidentally, T (°C) can be calculated by the following formula.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
１．０ｓ以上放冷
上記条件で圧延した後、放冷することにより、一部ひずみを開放し、続く冷却中のひずみ誘起析出や転位上析出を抑制し、Ｔｉ含有析出物やＮｂ含有析出物を低減することができる。このような効果を得るには圧延後の放冷時間を１．０ｓ以上とする必要がある。前記放冷時間は、好ましくは１．５ｓ以上とし、より好ましくは２．０ｓ以上とし、さらに好ましくは２．２ｓ以上とする。前記放冷時間の上限は特に限定されないが、前記放冷時間が５．０ｓ以下であると、その後の熱延制御をし易くなるため前記放冷時間は５．０ｓ以下が好ましい。なお、放冷とは、注水等による積極的な冷却（加速冷却）を行わずに大気中に暴露（空冷）することを意味する。なお、本発明において、熱間圧延は、粗圧延と仕上げ圧延を含み、前記圧延後の放冷時間は、熱間圧延後、すなわち、仕上げ圧延後の放冷時間である。 Cooling for 1.0 s or more By cooling after rolling under the above conditions, partial strain is released, strain-induced precipitation and dislocation precipitation during subsequent cooling are suppressed, and Ti-containing precipitates and Nb-containing precipitates can be reduced. To obtain such an effect, it is necessary to set the cooling time after rolling to 1.0 s or more. The cooling time is preferably 1.5 s or more, more preferably 2.0 s or more, and even more preferably 2.2 s or more. There is no particular limit to the upper limit of the cooling time, but if the cooling time is 5.0 s or less, it becomes easier to control the subsequent hot rolling, so the cooling time is preferably 5.0 s or less. Note that cooling means exposure to the atmosphere (air cooling) without active cooling (accelerated cooling) by water injection or the like. Note that in the present invention, hot rolling includes rough rolling and finish rolling, and the cooling time after rolling is the cooling time after hot rolling, i.e., after finish rolling.
上記条件で圧延した後、放冷することにより、一部ひずみを開放し、続く冷却中のひずみ誘起析出や転位上析出を抑制し、Ｔｉ含有析出物やＮｂ含有析出物を低減することができる。このような効果を得るには圧延後の放冷時間を１．０ｓ以上とする必要がある。前記放冷時間は、好ましくは１．５ｓ以上とし、より好ましくは２．０ｓ以上とし、さらに好ましくは２．２ｓ以上とする。前記放冷時間の上限は特に限定されないが、前記放冷時間が５．０ｓ以下であると、その後の熱延制御をし易くなるため前記放冷時間は５．０ｓ以下が好ましい。なお、放冷とは、注水等による積極的な冷却（加速冷却）を行わずに大気中に暴露（空冷）することを意味する。なお、本発明において、熱間圧延は、粗圧延と仕上げ圧延を含み、前記圧延後の放冷時間は、熱間圧延後、すなわち、仕上げ圧延後の放冷時間である。 Cooling for 1.0 s or more By cooling after rolling under the above conditions, partial strain is released, strain-induced precipitation and dislocation precipitation during subsequent cooling are suppressed, and Ti-containing precipitates and Nb-containing precipitates can be reduced. To obtain such an effect, it is necessary to set the cooling time after rolling to 1.0 s or more. The cooling time is preferably 1.5 s or more, more preferably 2.0 s or more, and even more preferably 2.2 s or more. There is no particular limit to the upper limit of the cooling time, but if the cooling time is 5.0 s or less, it becomes easier to control the subsequent hot rolling, so the cooling time is preferably 5.0 s or less. Note that cooling means exposure to the atmosphere (air cooling) without active cooling (accelerated cooling) by water injection or the like. Note that in the present invention, hot rolling includes rough rolling and finish rolling, and the cooling time after rolling is the cooling time after hot rolling, i.e., after finish rolling.
５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却
上記放冷後、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却する。５５０℃までの平均冷却速度が５０℃／ｓ未満では、フェライト、上部ベイナイト、Ｔｉ含有析出物、Ｎｂ含有析出物等の過剰な生成や表層領域の結晶方位の形成阻害を招く。その結果、本発明の相組織や、析出物、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、前記放冷後の冷却開始温度から５５０℃までの温度域の平均冷却速度は５０℃／ｓ以上とする。前記平均冷却速度は、好ましくは７０℃／ｓ以上とする。前記平均冷却速度の上限は特に限定されないが、前記平均冷却速度が５００℃／ｓ以上では鋼板形状の劣化を招く場合があるため、前記平均冷却速度は５００℃／ｓ未満が好ましく、２００℃／ｓ未満がより好ましい。 Cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C After the above cooling, cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C. If the average cooling rate up to 550°C is less than 50°C/s, excessive generation of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc., and formation of the crystal orientation in the surface layer region will be caused. As a result, the phase structure of the present invention, the precipitates, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 will not be obtained. Therefore, the average cooling rate in the temperature range from the cooling start temperature to 550°C after the above cooling is set to 50°C/s or more. The average cooling rate is preferably set to 70°C/s or more. There is no particular limit to the upper limit of the average cooling rate, but since an average cooling rate of 500°C/s or more may cause deterioration of the steel sheet shape, the average cooling rate is preferably less than 500°C/s, and more preferably less than 200°C/s.
上記放冷後、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却する。５５０℃までの平均冷却速度が５０℃／ｓ未満では、フェライト、上部ベイナイト、Ｔｉ含有析出物、Ｎｂ含有析出物等の過剰な生成や表層領域の結晶方位の形成阻害を招く。その結果、本発明の相組織や、析出物、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。したがって、前記放冷後の冷却開始温度から５５０℃までの温度域の平均冷却速度は５０℃／ｓ以上とする。前記平均冷却速度は、好ましくは７０℃／ｓ以上とする。前記平均冷却速度の上限は特に限定されないが、前記平均冷却速度が５００℃／ｓ以上では鋼板形状の劣化を招く場合があるため、前記平均冷却速度は５００℃／ｓ未満が好ましく、２００℃／ｓ未満がより好ましい。 Cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C After the above cooling, cooling at an average cooling rate of 50°C/s or more in the temperature range up to 550°C. If the average cooling rate up to 550°C is less than 50°C/s, excessive generation of ferrite, upper bainite, Ti-containing precipitates, Nb-containing precipitates, etc., and formation of the crystal orientation in the surface layer region will be caused. As a result, the phase structure of the present invention, the precipitates, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 will not be obtained. Therefore, the average cooling rate in the temperature range from the cooling start temperature to 550°C after the above cooling is set to 50°C/s or more. The average cooling rate is preferably set to 70°C/s or more. There is no particular limit to the upper limit of the average cooling rate, but since an average cooling rate of 500°C/s or more may cause deterioration of the steel sheet shape, the average cooling rate is preferably less than 500°C/s, and more preferably less than 200°C/s.
５５０℃に達してから急冷を開始するまでの時間：０．５～４．０ｓ
５５０℃に達してから急冷（後述の冷却速度２００℃／ｓ以上での急冷）を開始するまでの間に一定時間を確保する（一定時間の間隔をおく）ことで、表層近傍での中温域のベイナイトを形成させることができる。その結果、本発明の表層領域における｛１１０｝＜１１１＞方位の極密度を得ることができる。５５０℃に達してから急冷を開始するまでの時間が０．５ｓ未満ではこのような効果が十分得られず、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。一方、前記時間が４．０ｓを超えると上部ベイナイトが過剰に生成して本発明の相組織が得られなくなる。したがって、５５０℃から急冷を開始するまでの間の時間を０．５～４．０ｓとする。前記時間は、好ましくは０．７ｓ以上とする。また、前記時間は、好ましくは２．０ｓ以下とし、より好ましくは１．６ｓ以下とする。 Time from reaching 550°C to starting quenching: 0.5 to 4.0 s
By ensuring a certain time (leaving a certain time interval) between reaching 550°C and starting quenching (quenching at a cooling rate of 200°C/s or more, which will be described later), bainite can be formed in the medium temperature region near the surface layer. As a result, the pole density of the {110}<111> orientation in the surface layer region of the present invention can be obtained. If the time from reaching 550°C to starting quenching is less than 0.5 s, such an effect cannot be obtained sufficiently, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. On the other hand, if the time exceeds 4.0 s, upper bainite is excessively generated and the phase structure of the present invention cannot be obtained. Therefore, the time from 550°C to starting quenching is set to 0.5 to 4.0 s. The time is preferably 0.7 s or more. The time is preferably 2.0 s or less, more preferably 1.6 s or less.
５５０℃に達してから急冷（後述の冷却速度２００℃／ｓ以上での急冷）を開始するまでの間に一定時間を確保する（一定時間の間隔をおく）ことで、表層近傍での中温域のベイナイトを形成させることができる。その結果、本発明の表層領域における｛１１０｝＜１１１＞方位の極密度を得ることができる。５５０℃に達してから急冷を開始するまでの時間が０．５ｓ未満ではこのような効果が十分得られず、表層領域における｛１１０｝＜１１１＞方位の極密度の１．８～５．０の値が得られなくなる。一方、前記時間が４．０ｓを超えると上部ベイナイトが過剰に生成して本発明の相組織が得られなくなる。したがって、５５０℃から急冷を開始するまでの間の時間を０．５～４．０ｓとする。前記時間は、好ましくは０．７ｓ以上とする。また、前記時間は、好ましくは２．０ｓ以下とし、より好ましくは１．６ｓ以下とする。 Time from reaching 550°C to starting quenching: 0.5 to 4.0 s
By ensuring a certain time (leaving a certain time interval) between reaching 550°C and starting quenching (quenching at a cooling rate of 200°C/s or more, which will be described later), bainite can be formed in the medium temperature region near the surface layer. As a result, the pole density of the {110}<111> orientation in the surface layer region of the present invention can be obtained. If the time from reaching 550°C to starting quenching is less than 0.5 s, such an effect cannot be obtained sufficiently, and the pole density of the {110}<111> orientation in the surface layer region of 1.8 to 5.0 cannot be obtained. On the other hand, if the time exceeds 4.0 s, upper bainite is excessively generated and the phase structure of the present invention cannot be obtained. Therefore, the time from 550°C to starting quenching is set to 0.5 to 4.0 s. The time is preferably 0.7 s or more. The time is preferably 2.0 s or less, more preferably 1.6 s or less.
１００～２５０℃の巻取り温度までの冷却速度：２００℃／ｓ以上
上記のように、５５０℃に達してから急冷を開始するまでの間に０．５～４．０ｓの時間をおいた後、急冷を開始する。１００～２５０℃の巻取り温度までの冷却（急冷）速度が２００℃／ｓ未満では、上部ベイナイトや残留γが過剰に生じたり、表層領域における｛１１０｝＜１１１＞方位の極密度の増加を招く。その結果、本発明の相組織や、表層領域における｛１１０｝＜１１１＞方位の極密度が得られなくなる。したがって、巻取り温度までの冷却速度は２００℃／ｓ以上とする。前記冷却速度は、好ましくは２５０℃／ｓ以上とする。なお、前記冷却速度の上限は、特に限定されないが、形状安定性などの観点から、前記冷却速度は１０００℃／以下が好ましく、５００℃／ｓ以下がより好ましい。 Cooling rate to coiling temperature of 100 to 250°C: 200°C/s or more As described above, quenching is started after a time of 0.5 to 4.0 s has elapsed between reaching 550°C and starting quenching. If the cooling (quenching) rate to coiling temperature of 100 to 250°C is less than 200°C/s, upper bainite and residual γ are excessively generated, and the pole density of the {110}<111> orientation in the surface layer region is increased. As a result, the phase structure of the present invention and the pole density of the {110}<111> orientation in the surface layer region cannot be obtained. Therefore, the cooling rate to the coiling temperature is set to 200°C/s or more. The cooling rate is preferably set to 250°C/s or more. The upper limit of the cooling rate is not particularly limited, but from the viewpoint of shape stability, the cooling rate is preferably 1000°C/s or less, and more preferably 500°C/s or less.
上記のように、５５０℃に達してから急冷を開始するまでの間に０．５～４．０ｓの時間をおいた後、急冷を開始する。１００～２５０℃の巻取り温度までの冷却（急冷）速度が２００℃／ｓ未満では、上部ベイナイトや残留γが過剰に生じたり、表層領域における｛１１０｝＜１１１＞方位の極密度の増加を招く。その結果、本発明の相組織や、表層領域における｛１１０｝＜１１１＞方位の極密度が得られなくなる。したがって、巻取り温度までの冷却速度は２００℃／ｓ以上とする。前記冷却速度は、好ましくは２５０℃／ｓ以上とする。なお、前記冷却速度の上限は、特に限定されないが、形状安定性などの観点から、前記冷却速度は１０００℃／以下が好ましく、５００℃／ｓ以下がより好ましい。 Cooling rate to coiling temperature of 100 to 250°C: 200°C/s or more As described above, quenching is started after a time of 0.5 to 4.0 s has elapsed between reaching 550°C and starting quenching. If the cooling (quenching) rate to coiling temperature of 100 to 250°C is less than 200°C/s, upper bainite and residual γ are excessively generated, and the pole density of the {110}<111> orientation in the surface layer region is increased. As a result, the phase structure of the present invention and the pole density of the {110}<111> orientation in the surface layer region cannot be obtained. Therefore, the cooling rate to the coiling temperature is set to 200°C/s or more. The cooling rate is preferably set to 250°C/s or more. The upper limit of the cooling rate is not particularly limited, but from the viewpoint of shape stability, the cooling rate is preferably 1000°C/s or less, and more preferably 500°C/s or less.
巻取り温度：１００～２５０℃
巻取り温度を１００～２５０℃の範囲に調整することでマルテンサイトや下部ベイナイトを適度に焼戻すとともに他相を排除し、本発明の組織を得ることができる。巻取り温度が１００℃未満では、このような効果が十分得られず、フレッシュマルテンサイトが過剰に生じて本発明の組織が得られなくなる。一方、巻取り温度が２５０℃を超えるとマルテンサイトや下部ベイナイトの焼戻しが顕著になるとともにフレッシュマルテンサイトや残留γが生成して、本発明の組織が得られなくなる。したがって、巻取り温度は１００～２５０℃とする。巻取り温度は、好ましくは１２０℃以上とする。また、巻取り温度は、好ましくは２２０℃以下とする。 Winding temperature: 100 to 250°C
By adjusting the coiling temperature to the range of 100 to 250°C, martensite and lower bainite are appropriately tempered while other phases are eliminated, and the structure of the present invention can be obtained. If the coiling temperature is less than 100°C, such effects cannot be sufficiently obtained, and fresh martensite is excessively generated, making it impossible to obtain the structure of the present invention. On the other hand, if the coiling temperature exceeds 250°C, tempering of martensite and lower bainite becomes significant, and fresh martensite and residual γ are generated, making it impossible to obtain the structure of the present invention. Therefore, the coiling temperature is set to 100 to 250°C. The coiling temperature is preferably 120°C or higher. Moreover, the coiling temperature is preferably 220°C or lower.
巻取り温度を１００～２５０℃の範囲に調整することでマルテンサイトや下部ベイナイトを適度に焼戻すとともに他相を排除し、本発明の組織を得ることができる。巻取り温度が１００℃未満では、このような効果が十分得られず、フレッシュマルテンサイトが過剰に生じて本発明の組織が得られなくなる。一方、巻取り温度が２５０℃を超えるとマルテンサイトや下部ベイナイトの焼戻しが顕著になるとともにフレッシュマルテンサイトや残留γが生成して、本発明の組織が得られなくなる。したがって、巻取り温度は１００～２５０℃とする。巻取り温度は、好ましくは１２０℃以上とする。また、巻取り温度は、好ましくは２２０℃以下とする。 Winding temperature: 100 to 250°C
By adjusting the coiling temperature to the range of 100 to 250°C, martensite and lower bainite are appropriately tempered while other phases are eliminated, and the structure of the present invention can be obtained. If the coiling temperature is less than 100°C, such effects cannot be sufficiently obtained, and fresh martensite is excessively generated, making it impossible to obtain the structure of the present invention. On the other hand, if the coiling temperature exceeds 250°C, tempering of martensite and lower bainite becomes significant, and fresh martensite and residual γ are generated, making it impossible to obtain the structure of the present invention. Therefore, the coiling temperature is set to 100 to 250°C. The coiling temperature is preferably 120°C or higher. Moreover, the coiling temperature is preferably 220°C or lower.
上記した製造方法の条件以外は特に限定しないが、以下のように適宜条件を調整して製造することが好ましい。例えば、仕上げ圧延は、加工性の低下を招く粗粒低減等の観点から４パス以上とすることが好ましい。また、熱間圧延後は形状矯正や表面素度の調整等のために調質圧延を施しても構わない。酸洗を行う場合は５０～１００℃の酸洗浴に複数回浸漬して行うことが好ましい。
There are no particular limitations other than the conditions of the manufacturing method described above, but it is preferable to manufacture by adjusting the conditions as appropriate as follows. For example, it is preferable to perform finish rolling four or more passes in order to reduce coarse grains that can cause a decrease in workability. In addition, after hot rolling, temper rolling may be performed to correct the shape and adjust the surface roughness. If pickling is performed, it is preferable to perform multiple immersions in a pickling bath at 50 to 100°C.
本発明の高強度熱延鋼板は、後加熱後の強度、靭性、耐遅れ破壊特性に優れる。ここで、後加熱の加熱温度としては、４００℃以上が挙げられる。また、後加熱の加熱温度の上限は、特に限定されないが、一例として、後加熱の加熱温度は１１５０℃以下が挙げられる。後加熱の加熱時間（前記加熱温度での保持時間）は、特に限定されないが、一例として、０ｓ超が挙げられる。また、前記加熱時間は、一例として、３６００ｓ以下が挙げられる。
The high-strength hot-rolled steel sheet of the present invention has excellent strength, toughness, and delayed fracture resistance after post-heating. Here, the heating temperature of the post-heating is 400°C or higher. The upper limit of the heating temperature of the post-heating is not particularly limited, but an example of the heating temperature of the post-heating is 1150°C or lower. The heating time of the post-heating (holding time at the heating temperature) is not particularly limited, but an example of the heating time is more than 0 seconds. The heating time of the post-heating is 3600 seconds or less, for example.
表１に示す成分組成の鋼を転炉により溶製し、スラブとした後、表２に示す条件でスラブの加熱および熱間圧延を行い、熱延鋼板（原板）を製造した。得られた熱延鋼板を用いて、以下の試験方法に従い、組織観察、固溶Ｔｉ、固溶Ｎｂ、Ｔｉ含有析出物およびＮｂ含有析出物の分析、引張特性の評価を行った。さらに、前記熱延鋼板に表２に示す後加熱を施し、後加熱後の熱延鋼板を用いて、以下の試験方法に従い、硬さ、靭性および耐遅れ破壊特性の評価を行った。後加熱温度は伸びフランジ性の向上が認められるようになる４００℃以上とし、後加熱時間は生産性の観点から３６００ｓ以下の条件で行った。
Steel having the composition shown in Table 1 was melted in a converter and formed into a slab, which was then heated and hot rolled under the conditions shown in Table 2 to produce hot-rolled steel sheet (base sheet). The resulting hot-rolled steel sheet was used to observe the structure, analyze solute Ti, solute Nb, Ti-containing precipitates, and Nb-containing precipitates, and evaluate the tensile properties according to the following test methods. Furthermore, the hot-rolled steel sheet was post-heated as shown in Table 2, and the hardness, toughness, and delayed fracture resistance were evaluated according to the following test methods using the hot-rolled steel sheet after post-heating. The post-heating temperature was 400°C or higher, at which an improvement in stretch flangeability is observed, and the post-heating time was 3600 s or less from the viewpoint of productivity.
組織観察
マルテンサイトおよび下部ベイナイトの面積率とは、観察面積に占める各組織の面積の割合のことである。マルテンサイトの面積率は、得られた熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨後、３％ナイタールで腐食し、板厚１／４位置をＳＥＭ（走査型電子顕微鏡）で１５００倍の倍率でそれぞれ３視野撮影した。得られた２次電子像の画像データからＭｅｄｉａ Ｃｙｂｅｒｎｅｔｉｃｓ社製のＩｍａｇｅ－Ｐｒｏを用いて各組織の面積率を求め、３視野の平均面積率を各組織の面積率とした。組織の判定は一般的な分類により行って構わないが、例えば以下のように判定できる。画像データにおいて、下部ベイナイトは方位のそろった炭化物を含む黒または暗灰色または灰色または明灰色として区別される。マルテンサイトは規則的であるが複数の方位の炭化物を含む黒～明灰色の組織である。あるいは炭化物を含まない白色または明灰色として観察される。残留オーステナイトは炭化物を含まない白または明灰色として観察される。マルテンサイトの一部と残留オーステナイトは区別できない場合があるため、残留オーステナイトは後述する方法にて求め、ＳＥＭ像から求めたマルテンサイトと残留オーステナイトの合計面積率から除してマルテンサイトの面積率を求めた。なお、焼戻しの程度が強い組織ほど、素地は黒が強いコントラストの像となるため、上記素地の色は目安であり、本発明では炭化物の量や組織形態等を総合して判断し、後述の組織を含め、特徴が近いいずれかの組織に分類した。炭化物は白色の点状または線状である。また、上記以外の組織として、フェライトは黒または暗灰色で内部に炭化物やラス等の下部組織を有さない組織であり、パーライトは黒色と白色の層状または部分的に途切れた層状組織として区別できる。また、上部ベイナイトは、黒または暗灰色で内部に炭化物やラス等の下部組織を有する組織として区別できる。残留γ量は、次のように求める。熱延鋼板を板厚の１／４＋０．１ｍｍまで研削後、化学研磨によりさらに０．１ｍｍ研磨した面を測定面とする。前記測定面について、Ｘ線回折装置でＭｏのＫα１線を用い、ｆｃｃ鉄（オーステナイト）の（２００）面、（２２０）面、（３１１）面と、ｂｃｃ鉄（フェライト）の（２００）面、（２１１）面、（２２０）面の積分反射強度を測定する。そして、ｂｃｃ鉄の各面からの積分反射強度に対するｆｃｃ鉄の各面からの積分反射強度の強度比から体積率を求め、これを残留γ量とした。 Structure observation The area ratio of martensite and lower bainite refers to the ratio of the area of each structure to the observed area. The area ratio of martensite was measured by cutting a sample from the obtained hot-rolled steel sheet, polishing the plate thickness cross section parallel to the rolling direction, corroding it with 3% nital, and photographing the plate thickness 1/4 position at a magnification of 1500 times with a SEM (scanning electron microscope) in three fields of view. The area ratio of each structure was calculated from the image data of the obtained secondary electron image using Image-Pro manufactured by Media Cybernetics, and the average area ratio of the three fields of view was taken as the area ratio of each structure. The structure may be determined by a general classification, but can be determined, for example, as follows. In the image data, lower bainite is distinguished as black or dark gray, gray, or light gray containing oriented carbides. Martensite is a structure of black to light gray that is regular but contains carbides of multiple orientations. Alternatively, it is observed as white or light gray without carbides. The retained austenite is observed as white or light gray without containing carbides. Since it may be difficult to distinguish between a part of martensite and the retained austenite, the retained austenite was obtained by the method described below, and the area ratio of the martensite was obtained by subtracting it from the total area ratio of the martensite and the retained austenite obtained from the SEM image. The stronger the degree of tempering, the stronger the contrast of the base material is in black, so the color of the base material is a guide, and in the present invention, the amount of carbides, the structure form, etc. are comprehensively judged, and the structures are classified into one of the structures with similar characteristics, including the structures described below. The carbides are white dots or lines. In addition to the above structures, ferrite is a structure that is black or dark gray and does not have a substructure such as carbides or laths inside, and pearlite can be distinguished as a black and white layered or partially interrupted layered structure. In addition, upper bainite can be distinguished as a structure that is black or dark gray and has a substructure such as carbides or laths inside. The amount of retained γ is obtained as follows. The hot-rolled steel sheet was ground to 1/4+0.1 mm of the sheet thickness, and then chemically polished to a further 0.1 mm to obtain the measurement surface. The measurement surface was measured using an X-ray diffraction apparatus with Mo Kα1 radiation to measure the integral reflection intensities of the (200), (220), and (311) surfaces of fcc iron (austenite) and the (200), (211), and (220) surfaces of bcc iron (ferrite). The volume fraction was then calculated from the intensity ratio of the integral reflection intensity from each surface of the fcc iron to the integral reflection intensity from each surface of the bcc iron, and this was taken as the amount of residual γ.
マルテンサイトおよび下部ベイナイトの面積率とは、観察面積に占める各組織の面積の割合のことである。マルテンサイトの面積率は、得られた熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨後、３％ナイタールで腐食し、板厚１／４位置をＳＥＭ（走査型電子顕微鏡）で１５００倍の倍率でそれぞれ３視野撮影した。得られた２次電子像の画像データからＭｅｄｉａ Ｃｙｂｅｒｎｅｔｉｃｓ社製のＩｍａｇｅ－Ｐｒｏを用いて各組織の面積率を求め、３視野の平均面積率を各組織の面積率とした。組織の判定は一般的な分類により行って構わないが、例えば以下のように判定できる。画像データにおいて、下部ベイナイトは方位のそろった炭化物を含む黒または暗灰色または灰色または明灰色として区別される。マルテンサイトは規則的であるが複数の方位の炭化物を含む黒～明灰色の組織である。あるいは炭化物を含まない白色または明灰色として観察される。残留オーステナイトは炭化物を含まない白または明灰色として観察される。マルテンサイトの一部と残留オーステナイトは区別できない場合があるため、残留オーステナイトは後述する方法にて求め、ＳＥＭ像から求めたマルテンサイトと残留オーステナイトの合計面積率から除してマルテンサイトの面積率を求めた。なお、焼戻しの程度が強い組織ほど、素地は黒が強いコントラストの像となるため、上記素地の色は目安であり、本発明では炭化物の量や組織形態等を総合して判断し、後述の組織を含め、特徴が近いいずれかの組織に分類した。炭化物は白色の点状または線状である。また、上記以外の組織として、フェライトは黒または暗灰色で内部に炭化物やラス等の下部組織を有さない組織であり、パーライトは黒色と白色の層状または部分的に途切れた層状組織として区別できる。また、上部ベイナイトは、黒または暗灰色で内部に炭化物やラス等の下部組織を有する組織として区別できる。残留γ量は、次のように求める。熱延鋼板を板厚の１／４＋０．１ｍｍまで研削後、化学研磨によりさらに０．１ｍｍ研磨した面を測定面とする。前記測定面について、Ｘ線回折装置でＭｏのＫα１線を用い、ｆｃｃ鉄（オーステナイト）の（２００）面、（２２０）面、（３１１）面と、ｂｃｃ鉄（フェライト）の（２００）面、（２１１）面、（２２０）面の積分反射強度を測定する。そして、ｂｃｃ鉄の各面からの積分反射強度に対するｆｃｃ鉄の各面からの積分反射強度の強度比から体積率を求め、これを残留γ量とした。 Structure observation The area ratio of martensite and lower bainite refers to the ratio of the area of each structure to the observed area. The area ratio of martensite was measured by cutting a sample from the obtained hot-rolled steel sheet, polishing the plate thickness cross section parallel to the rolling direction, corroding it with 3% nital, and photographing the plate thickness 1/4 position at a magnification of 1500 times with a SEM (scanning electron microscope) in three fields of view. The area ratio of each structure was calculated from the image data of the obtained secondary electron image using Image-Pro manufactured by Media Cybernetics, and the average area ratio of the three fields of view was taken as the area ratio of each structure. The structure may be determined by a general classification, but can be determined, for example, as follows. In the image data, lower bainite is distinguished as black or dark gray, gray, or light gray containing oriented carbides. Martensite is a structure of black to light gray that is regular but contains carbides of multiple orientations. Alternatively, it is observed as white or light gray without carbides. The retained austenite is observed as white or light gray without containing carbides. Since it may be difficult to distinguish between a part of martensite and the retained austenite, the retained austenite was obtained by the method described below, and the area ratio of the martensite was obtained by subtracting it from the total area ratio of the martensite and the retained austenite obtained from the SEM image. The stronger the degree of tempering, the stronger the contrast of the base material is in black, so the color of the base material is a guide, and in the present invention, the amount of carbides, the structure form, etc. are comprehensively judged, and the structures are classified into one of the structures with similar characteristics, including the structures described below. The carbides are white dots or lines. In addition to the above structures, ferrite is a structure that is black or dark gray and does not have a substructure such as carbides or laths inside, and pearlite can be distinguished as a black and white layered or partially interrupted layered structure. In addition, upper bainite can be distinguished as a structure that is black or dark gray and has a substructure such as carbides or laths inside. The amount of retained γ is obtained as follows. The hot-rolled steel sheet was ground to 1/4+0.1 mm of the sheet thickness, and then chemically polished to a further 0.1 mm to obtain the measurement surface. The measurement surface was measured using an X-ray diffraction apparatus with Mo Kα1 radiation to measure the integral reflection intensities of the (200), (220), and (311) surfaces of fcc iron (austenite) and the (200), (211), and (220) surfaces of bcc iron (ferrite). The volume fraction was then calculated from the intensity ratio of the integral reflection intensity from each surface of the fcc iron to the integral reflection intensity from each surface of the bcc iron, and this was taken as the amount of residual γ.
得られた各組織の面積率を用いて５０％以上となる主相を構成する組織とその他の組織を表３に示す。なお、表３中のＭはマルテンサイト、ＬＢは下部ベイナイト、γは残留オーステナイト、Ｏはその他の相を意味する。その他の相には、フェライト、パーライト、上部ベイナイトの１種または２種以上が含まれる。
Table 3 shows the structures constituting the main phase and other structures that account for 50% or more of the area ratio of each structure obtained. In Table 3, M means martensite, LB means lower bainite, γ means retained austenite, and O means other phases. The other phases include one or more of ferrite, pearlite, and upper bainite.
表面から板厚中央方向に１００μｍまでの表層領域における｛１１０｝＜１１１＞方位の極密度
得られた熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨し、さらに電解研磨によりひずみを除去した後、表面から板厚中央方向に１００μｍまでの表層領域について、ＥＢＳＤ（電子線後方散乱回折）法により結晶方位データを取得した。測定領域を１００μｍ×１００μｍ、加速電圧を３０ｋＶ、ステップサイズを１００ｎｍとして各サンプルについて３視野測定した。取得データの解析にはＴＳＬソリューション社製のＯＩＭＡｎａｌｙｓｉｓ Ｖｅｒ．７．３．０を用いた。取得データについて、Ｃｏｎｆｉｄｅｎｃｅ Ｉｎｄｅｘ（ＣＩ）値が０．１以下のデータをカットした後、φ１＝３０～４０゜、φ２＝４５゜、Φ＝８５～９０゜とし、かつそれぞれのＲｅｓｏｌｕｔｉｏｎを５゜としたＯｒｉｅｎｔａｔｉｏｎＤｉｓｔｒｉｂｕｔｉｏｎ Ｆｕｎｃｔｉｏｎ（ＯＤＦ）の計算を行い、その領域の極密度を算出し、その平均値を各視野の｛１１０｝＜１１１＞方位の極密度とした。そして、各サンプルの３視野の極密度の平均値を、各サンプルの極密度とした。 Pole density of {110}<111> orientation in the surface region from the surface to 100 μm in the thickness center direction Samples were cut out from the obtained hot-rolled steel sheet, the thickness cross section parallel to the rolling direction was polished, and strain was removed by electrolytic polishing. After that, crystal orientation data was obtained for the surface region from the surface to 100 μm in the thickness center direction by EBSD (electron backscatter diffraction). Each sample was measured in three fields of view with a measurement area of 100 μm x 100 μm, an acceleration voltage of 30 kV, and a step size of 100 nm. OIM Analysis Ver. 7.3.0 manufactured by TSL Solutions was used to analyze the obtained data. After cutting out data with a Confidence Index (CI) value of 0.1 or less from the acquired data, the Orientation Distribution Function (ODF) was calculated with φ1 = 30 to 40°, φ2 = 45°, Φ = 85 to 90°, and each Resolution of 5°, and the pole density of the region was calculated, and the average value was taken as the pole density of the {110}<111> orientation of each field of view. The average value of the pole densities of the three fields of view of each sample was taken as the pole density of each sample.
得られた熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨し、さらに電解研磨によりひずみを除去した後、表面から板厚中央方向に１００μｍまでの表層領域について、ＥＢＳＤ（電子線後方散乱回折）法により結晶方位データを取得した。測定領域を１００μｍ×１００μｍ、加速電圧を３０ｋＶ、ステップサイズを１００ｎｍとして各サンプルについて３視野測定した。取得データの解析にはＴＳＬソリューション社製のＯＩＭＡｎａｌｙｓｉｓ Ｖｅｒ．７．３．０を用いた。取得データについて、Ｃｏｎｆｉｄｅｎｃｅ Ｉｎｄｅｘ（ＣＩ）値が０．１以下のデータをカットした後、φ１＝３０～４０゜、φ２＝４５゜、Φ＝８５～９０゜とし、かつそれぞれのＲｅｓｏｌｕｔｉｏｎを５゜としたＯｒｉｅｎｔａｔｉｏｎＤｉｓｔｒｉｂｕｔｉｏｎ Ｆｕｎｃｔｉｏｎ（ＯＤＦ）の計算を行い、その領域の極密度を算出し、その平均値を各視野の｛１１０｝＜１１１＞方位の極密度とした。そして、各サンプルの３視野の極密度の平均値を、各サンプルの極密度とした。 Pole density of {110}<111> orientation in the surface region from the surface to 100 μm in the thickness center direction Samples were cut out from the obtained hot-rolled steel sheet, the thickness cross section parallel to the rolling direction was polished, and strain was removed by electrolytic polishing. After that, crystal orientation data was obtained for the surface region from the surface to 100 μm in the thickness center direction by EBSD (electron backscatter diffraction). Each sample was measured in three fields of view with a measurement area of 100 μm x 100 μm, an acceleration voltage of 30 kV, and a step size of 100 nm. OIM Analysis Ver. 7.3.0 manufactured by TSL Solutions was used to analyze the obtained data. After cutting out data with a Confidence Index (CI) value of 0.1 or less from the acquired data, the Orientation Distribution Function (ODF) was calculated with φ1 = 30 to 40°, φ2 = 45°, Φ = 85 to 90°, and each Resolution of 5°, and the pole density of the region was calculated, and the average value was taken as the pole density of the {110}<111> orientation of each field of view. The average value of the pole densities of the three fields of view of each sample was taken as the pole density of each sample.
固溶Ｔｉ、固溶Ｎｂ、Ｔｉ含有析出物、Ｎｂ含有析出物分析
得られた熱延鋼板より幅３０ｍｍ、長さ３０ｍｍの試験片を採取し、非水溶媒系電解液（１０％ＡＡ系電解液：１０ｖｏｌ％アセチルアセトン－１ｍａｓｓ％塩化テトラメチルアンモニウム－メタノール）中で定電流電解を行った。電流密度は２０ｍＡ/ｃｍ２とし、電解量は約０．２ｇとした。電解後の電解液を分析溶液とし、ＩＣＰ質量分析法を用いてＴｉ、Ｎｂおよび比較元素としてＦｅの液中濃度（質量％）を測定した。得られた濃度を基に、Ｆｅに対するＴｉ、Ｎｂの濃度比を算出し、さらに、試験片中のＦｅの含有量（質量％）を乗じることで、固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）とした。なお、試験片中のＦｅの含有量（質量％）は、Ｆｅ以外の成分含有量の合計（質量％）を１００質量％から差し引くことで求めた。得られた固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）を用いて、含有Ｔｉ量（質量％）および含有Ｎｂ量（質量％）の合計に対する固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）の合計の比率を算出した。一方、電解した後の、表面に析出物が付着している試験片を、電解液から取り出して、ヘキサメタリン酸ナトリウム水溶液（５００ｍｇ／Ｌ）（以下、ＳＨＭＰ水溶液と称す）中に浸漬した。そして、超音波振動を付与して、析出物を試験片から剥離しＳＨＭＰ水溶液中に抽出した。ついで、析出物を含むＳＨＭＰ水溶液を、孔径１００ｎｍのフィルタを用いてろ過し、次いで１００ｎｍフィルタに捕集された析出物を酸分解し、分解液に対してＩＣＰ発光分光分析装置を用いて分析し、分解液中のＴｉおよびＮｂの絶対値を測定した。得られたＴｉおよびＮｂの絶対値を電解質量で除し、粒径１００ｎｍ以上の析出物に含まれるＴｉ量およびＮｂ量（試験片の全組成を１００質量％とした場合の質量％）を得た。次に、得られたＴｉ量（質量％）およびＮｂ量（質量％）の合計を、試験片中の含有Ｔｉ量（質量％）および含有Ｎｂ量（質量％）の合計で除し、粒径１００ｎｍ以上のＴｉを含む析出物として存在しているＴｉ量（質量％）と粒径１００ｎｍ以上のＮｂを含む析出物量として存在しているＮｂ量（質量％）の合計とした。なお、電解質量は、析出物剥離後の試験片に対して質量を測定し、電解前の試験片質量から差し引くことで求めた。 Analysis of dissolved Ti, dissolved Nb, Ti-containing precipitates, and Nb-containing precipitates A test piece with a width of 30 mm and a length of 30 mm was taken from the obtained hot-rolled steel sheet, and constant current electrolysis was performed in a non-aqueous solvent-based electrolyte (10% AA-based electrolyte: 10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol). The current density was 20 mA/cm 2 , and the amount of electrolysis was about 0.2 g. The electrolyte after electrolysis was used as an analysis solution, and the concentrations (mass%) of Ti, Nb, and Fe as a comparative element were measured using ICP mass spectrometry. Based on the obtained concentration, the concentration ratios of Ti and Nb to Fe were calculated, and further multiplied by the content (mass%) of Fe in the test piece to obtain the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%). The content (mass%) of Fe in the test piece was obtained by subtracting the total content (mass%) of components other than Fe from 100 mass%. Using the obtained dissolved Ti amount (mass%) and dissolved Nb amount (mass%), the ratio of the total dissolved Ti amount (mass%) and dissolved Nb amount (mass%) to the total contained Ti amount (mass%) and contained Nb amount (mass%) was calculated. On the other hand, the test piece having the precipitate attached to the surface after electrolysis was taken out from the electrolytic solution and immersed in an aqueous solution of sodium hexametaphosphate (500 mg/L) (hereinafter referred to as an aqueous SHMP solution). Then, ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution. Next, the aqueous SHMP solution containing the precipitate was filtered using a filter with a pore size of 100 nm, and the precipitate collected on the 100 nm filter was decomposed with acid, and the decomposition solution was analyzed using an ICP emission spectrometer, and the absolute values of Ti and Nb in the decomposition solution were measured. The absolute values of Ti and Nb obtained were divided by the amount of electrolyte to obtain the amount of Ti and the amount of Nb (mass %) contained in the precipitates having a particle size of 100 nm or more when the total composition of the test piece was taken as 100 mass %). Next, the total of the obtained amount of Ti (mass %) and amount of Nb (mass %) was divided by the total amount of Ti (mass %) and amount of Nb (mass %) contained in the test piece to obtain the total amount of Ti (mass %) present as precipitates containing Ti having a particle size of 100 nm or more and the amount of Nb (mass %) present as precipitates containing Nb having a particle size of 100 nm or more. The amount of electrolyte was obtained by measuring the mass of the test piece after the precipitates were peeled off and subtracting it from the mass of the test piece before electrolysis.
得られた熱延鋼板より幅３０ｍｍ、長さ３０ｍｍの試験片を採取し、非水溶媒系電解液（１０％ＡＡ系電解液：１０ｖｏｌ％アセチルアセトン－１ｍａｓｓ％塩化テトラメチルアンモニウム－メタノール）中で定電流電解を行った。電流密度は２０ｍＡ/ｃｍ２とし、電解量は約０．２ｇとした。電解後の電解液を分析溶液とし、ＩＣＰ質量分析法を用いてＴｉ、Ｎｂおよび比較元素としてＦｅの液中濃度（質量％）を測定した。得られた濃度を基に、Ｆｅに対するＴｉ、Ｎｂの濃度比を算出し、さらに、試験片中のＦｅの含有量（質量％）を乗じることで、固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）とした。なお、試験片中のＦｅの含有量（質量％）は、Ｆｅ以外の成分含有量の合計（質量％）を１００質量％から差し引くことで求めた。得られた固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）を用いて、含有Ｔｉ量（質量％）および含有Ｎｂ量（質量％）の合計に対する固溶Ｔｉ量（質量％）および固溶Ｎｂ量（質量％）の合計の比率を算出した。一方、電解した後の、表面に析出物が付着している試験片を、電解液から取り出して、ヘキサメタリン酸ナトリウム水溶液（５００ｍｇ／Ｌ）（以下、ＳＨＭＰ水溶液と称す）中に浸漬した。そして、超音波振動を付与して、析出物を試験片から剥離しＳＨＭＰ水溶液中に抽出した。ついで、析出物を含むＳＨＭＰ水溶液を、孔径１００ｎｍのフィルタを用いてろ過し、次いで１００ｎｍフィルタに捕集された析出物を酸分解し、分解液に対してＩＣＰ発光分光分析装置を用いて分析し、分解液中のＴｉおよびＮｂの絶対値を測定した。得られたＴｉおよびＮｂの絶対値を電解質量で除し、粒径１００ｎｍ以上の析出物に含まれるＴｉ量およびＮｂ量（試験片の全組成を１００質量％とした場合の質量％）を得た。次に、得られたＴｉ量（質量％）およびＮｂ量（質量％）の合計を、試験片中の含有Ｔｉ量（質量％）および含有Ｎｂ量（質量％）の合計で除し、粒径１００ｎｍ以上のＴｉを含む析出物として存在しているＴｉ量（質量％）と粒径１００ｎｍ以上のＮｂを含む析出物量として存在しているＮｂ量（質量％）の合計とした。なお、電解質量は、析出物剥離後の試験片に対して質量を測定し、電解前の試験片質量から差し引くことで求めた。 Analysis of dissolved Ti, dissolved Nb, Ti-containing precipitates, and Nb-containing precipitates A test piece with a width of 30 mm and a length of 30 mm was taken from the obtained hot-rolled steel sheet, and constant current electrolysis was performed in a non-aqueous solvent-based electrolyte (10% AA-based electrolyte: 10 vol% acetylacetone-1 mass% tetramethylammonium chloride-methanol). The current density was 20 mA/cm 2 , and the amount of electrolysis was about 0.2 g. The electrolyte after electrolysis was used as an analysis solution, and the concentrations (mass%) of Ti, Nb, and Fe as a comparative element were measured using ICP mass spectrometry. Based on the obtained concentration, the concentration ratios of Ti and Nb to Fe were calculated, and further multiplied by the content (mass%) of Fe in the test piece to obtain the amount of dissolved Ti (mass%) and the amount of dissolved Nb (mass%). The content (mass%) of Fe in the test piece was obtained by subtracting the total content (mass%) of components other than Fe from 100 mass%. Using the obtained dissolved Ti amount (mass%) and dissolved Nb amount (mass%), the ratio of the total dissolved Ti amount (mass%) and dissolved Nb amount (mass%) to the total contained Ti amount (mass%) and contained Nb amount (mass%) was calculated. On the other hand, the test piece having the precipitate attached to the surface after electrolysis was taken out from the electrolytic solution and immersed in an aqueous solution of sodium hexametaphosphate (500 mg/L) (hereinafter referred to as an aqueous SHMP solution). Then, ultrasonic vibration was applied to peel off the precipitate from the test piece and extract it into the aqueous SHMP solution. Next, the aqueous SHMP solution containing the precipitate was filtered using a filter with a pore size of 100 nm, and the precipitate collected on the 100 nm filter was decomposed with acid, and the decomposition solution was analyzed using an ICP emission spectrometer, and the absolute values of Ti and Nb in the decomposition solution were measured. The absolute values of Ti and Nb obtained were divided by the amount of electrolyte to obtain the amount of Ti and the amount of Nb (mass %) contained in the precipitates having a particle size of 100 nm or more when the total composition of the test piece was taken as 100 mass %). Next, the total of the obtained amount of Ti (mass %) and amount of Nb (mass %) was divided by the total amount of Ti (mass %) and amount of Nb (mass %) contained in the test piece to obtain the total amount of Ti (mass %) present as precipitates containing Ti having a particle size of 100 nm or more and the amount of Nb (mass %) present as precipitates containing Nb having a particle size of 100 nm or more. The amount of electrolyte was obtained by measuring the mass of the test piece after the precipitates were peeled off and subtracting it from the mass of the test piece before electrolysis.
引張試験
得られた熱延鋼板より、圧延方向に対して平行方向にＪＩＳ５号引張試験片（ＪＩＳ Ｚ ２２４１：２０１１）を採取し、歪速度が１０－３／ｓとするＪＩＳ Ｚ ２２４１：２０１１の規定に準拠した引張試験を行い、ＴＳを求めた。なお、本発明では、ＴＳは１１８０ＭＰａ以上を合格とした。 Tensile Test JIS No. 5 tensile test pieces (JIS Z 2241:2011) were taken from the obtained hot-rolled steel sheets in the direction parallel to the rolling direction, and a tensile test was carried out in accordance with the provisions of JIS Z 2241:2011 at a strain rate of 10 −3 /s to determine TS. In the present invention, a TS of 1180 MPa or more was considered to be acceptable.
得られた熱延鋼板より、圧延方向に対して平行方向にＪＩＳ５号引張試験片（ＪＩＳ Ｚ ２２４１：２０１１）を採取し、歪速度が１０－３／ｓとするＪＩＳ Ｚ ２２４１：２０１１の規定に準拠した引張試験を行い、ＴＳを求めた。なお、本発明では、ＴＳは１１８０ＭＰａ以上を合格とした。 Tensile Test JIS No. 5 tensile test pieces (JIS Z 2241:2011) were taken from the obtained hot-rolled steel sheets in the direction parallel to the rolling direction, and a tensile test was carried out in accordance with the provisions of JIS Z 2241:2011 at a strain rate of 10 −3 /s to determine TS. In the present invention, a TS of 1180 MPa or more was considered to be acceptable.
ビッカース硬さ試験
得られた熱延鋼板および後加熱後の熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨後、板厚１／４位置において、荷重を５ｋｇ、測定点数を５点として、ビッカース硬さ試験を行い、その平均（算術平均）を鋼板のビッカース硬さとした。後加熱前後の硬さの差（ΔＨＶ）が５０以下を後加熱後の強度に優れると判断し、合格とした。 Vickers hardness test Samples were cut out from the obtained hot-rolled steel sheet and the hot-rolled steel sheet after post-heating, and the cross section of the sheet thickness parallel to the rolling direction was polished. Then, a Vickers hardness test was performed at 1/4 of the sheet thickness position with a load of 5 kg and five measurement points, and the average (arithmetic mean) was taken as the Vickers hardness of the steel sheet. A difference in hardness (ΔHV) of 50 or less before and after post-heating was judged to be excellent in strength after post-heating and was considered to have passed the test.
得られた熱延鋼板および後加熱後の熱延鋼板よりサンプルを切り出し、圧延方向に平行な板厚断面を研磨後、板厚１／４位置において、荷重を５ｋｇ、測定点数を５点として、ビッカース硬さ試験を行い、その平均（算術平均）を鋼板のビッカース硬さとした。後加熱前後の硬さの差（ΔＨＶ）が５０以下を後加熱後の強度に優れると判断し、合格とした。 Vickers hardness test Samples were cut out from the obtained hot-rolled steel sheet and the hot-rolled steel sheet after post-heating, and the cross section of the sheet thickness parallel to the rolling direction was polished. Then, a Vickers hardness test was performed at 1/4 of the sheet thickness position with a load of 5 kg and five measurement points, and the average (arithmetic mean) was taken as the Vickers hardness of the steel sheet. A difference in hardness (ΔHV) of 50 or less before and after post-heating was judged to be excellent in strength after post-heating and was considered to have passed the test.
シャルピー衝撃試験
得られた熱延鋼板を後加熱処理した熱延鋼板より、幅が１０ｍｍ、長さが５５ｍｍの試験片を採取し、先端角４５゜、先端半径０．２５ｍｍ、深さ２ｍｍのＶノッチを入れたシャルピー衝撃試験片を作製した。そして、ＪＩＳ Ｚ ２２４２：２０１８に準拠して、シャルピー衝撃試験を－２０℃で５回行い、延性破面率を評価した。５回の延性破面率の平均値が５０％以上を後加熱後の靭性に優れると判断し、合格とした。なお、板厚は２．９ｍｍとし、ノッチ方向は圧延方向に平行とした。 Charpy impact test From the hot-rolled steel sheet obtained by post-heating the hot-rolled steel sheet, a test piece with a width of 10 mm and a length of 55 mm was taken, and a V-notch with a tip angle of 45°, a tip radius of 0.25 mm, and a depth of 2 mm was made to prepare a Charpy impact test piece. Then, in accordance with JIS Z 2242:2018, a Charpy impact test was performed five times at -20 ° C. to evaluate the ductile fracture rate. A test piece with an average ductile fracture rate of 50% or more after five tests was judged to have excellent toughness after post-heating and was passed. The plate thickness was 2.9 mm, and the notch direction was parallel to the rolling direction.
得られた熱延鋼板を後加熱処理した熱延鋼板より、幅が１０ｍｍ、長さが５５ｍｍの試験片を採取し、先端角４５゜、先端半径０．２５ｍｍ、深さ２ｍｍのＶノッチを入れたシャルピー衝撃試験片を作製した。そして、ＪＩＳ Ｚ ２２４２：２０１８に準拠して、シャルピー衝撃試験を－２０℃で５回行い、延性破面率を評価した。５回の延性破面率の平均値が５０％以上を後加熱後の靭性に優れると判断し、合格とした。なお、板厚は２．９ｍｍとし、ノッチ方向は圧延方向に平行とした。 Charpy impact test From the hot-rolled steel sheet obtained by post-heating the hot-rolled steel sheet, a test piece with a width of 10 mm and a length of 55 mm was taken, and a V-notch with a tip angle of 45°, a tip radius of 0.25 mm, and a depth of 2 mm was made to prepare a Charpy impact test piece. Then, in accordance with JIS Z 2242:2018, a Charpy impact test was performed five times at -20 ° C. to evaluate the ductile fracture rate. A test piece with an average ductile fracture rate of 50% or more after five tests was judged to have excellent toughness after post-heating and was passed. The plate thickness was 2.9 mm, and the notch direction was parallel to the rolling direction.
遅れ破壊試験
得られた熱延鋼板より、幅が３０ｍｍ、長さが１１０ｍｍの試験片を採取し、表２に示す後加熱処理を施し、試験片を得た。これを稜線が圧延方向と平行となるように１５ｍｍの曲げ半径で９０゜Ｖ曲げ加工を行い、スプリングバックで開いた分だけボルトで締め込み、ｐＨ３の塩酸に９６ｈｒ浸漬し、割れの有無を調査した。割れ発生が無かったものを後加熱後の耐遅れ破壊特性に優れると判断し、合格とした。なお、試験片の端面はシャー角１゜、クリアランス１０％でせん断で形成させ、バリを曲げ外側とした。なお、表３中の「遅れ破壊時間（ｈｒ）」は、試験片に割れが発生した時間を示す。ただし、前記「遅れ破壊時間（ｈｒ）」の「９６」は、上記９６ｈｒの浸漬後に、試験片に割れ発生が無かったことを示す。 Delayed fracture test From the obtained hot-rolled steel sheet, a test piece with a width of 30 mm and a length of 110 mm was taken, and the post-heating treatment shown in Table 2 was performed to obtain a test piece. This was subjected to 90° V-bending with a bending radius of 15 mm so that the ridge line was parallel to the rolling direction, and bolts were tightened by the amount of opening due to springback, and the test piece was immersed in hydrochloric acid of pH 3 for 96 hours to check for the presence or absence of cracks. Those that did not have cracks were judged to have excellent delayed fracture resistance after post-heating and were passed. The end faces of the test pieces were formed by shearing with a shear angle of 1° and a clearance of 10%, and burrs were bent on the outside. The "delayed fracture time (hr)" in Table 3 indicates the time when cracks occurred in the test piece. However, "96" in the "delayed fracture time (hr)" indicates that no cracks occurred in the test piece after the above 96-hr immersion.
得られた熱延鋼板より、幅が３０ｍｍ、長さが１１０ｍｍの試験片を採取し、表２に示す後加熱処理を施し、試験片を得た。これを稜線が圧延方向と平行となるように１５ｍｍの曲げ半径で９０゜Ｖ曲げ加工を行い、スプリングバックで開いた分だけボルトで締め込み、ｐＨ３の塩酸に９６ｈｒ浸漬し、割れの有無を調査した。割れ発生が無かったものを後加熱後の耐遅れ破壊特性に優れると判断し、合格とした。なお、試験片の端面はシャー角１゜、クリアランス１０％でせん断で形成させ、バリを曲げ外側とした。なお、表３中の「遅れ破壊時間（ｈｒ）」は、試験片に割れが発生した時間を示す。ただし、前記「遅れ破壊時間（ｈｒ）」の「９６」は、上記９６ｈｒの浸漬後に、試験片に割れ発生が無かったことを示す。 Delayed fracture test From the obtained hot-rolled steel sheet, a test piece with a width of 30 mm and a length of 110 mm was taken, and the post-heating treatment shown in Table 2 was performed to obtain a test piece. This was subjected to 90° V-bending with a bending radius of 15 mm so that the ridge line was parallel to the rolling direction, and bolts were tightened by the amount of opening due to springback, and the test piece was immersed in hydrochloric acid of pH 3 for 96 hours to check for the presence or absence of cracks. Those that did not have cracks were judged to have excellent delayed fracture resistance after post-heating and were passed. The end faces of the test pieces were formed by shearing with a shear angle of 1° and a clearance of 10%, and burrs were bent on the outside. The "delayed fracture time (hr)" in Table 3 indicates the time when cracks occurred in the test piece. However, "96" in the "delayed fracture time (hr)" indicates that no cracks occurred in the test piece after the above 96-hr immersion.
発明例は、いずれも１１８０ＭＰａ以上のＴＳを有し、さらに後加熱後の強度、靭性および耐遅れ破壊特性に優れる。一方、本発明の範囲を外れる比較例は、所望の強度（ＴＳ）を有しないか、後加熱後の所望の強度、靭性、耐遅れ破壊特性のいずれか一つ以上が得られていない。
All of the inventive examples have a TS of 1180 MPa or more, and are excellent in strength, toughness, and delayed fracture resistance after post-heating. On the other hand, the comparative examples that fall outside the scope of the present invention either do not have the desired strength (TS) or do not achieve one or more of the desired strength, toughness, and delayed fracture resistance after post-heating.
本発明によれば、ＴＳが１１８０ＭＰａ以上１６００ＭＰａ未満で、後加熱後の強度、靭性および耐遅れ破壊特性に優れる高強度熱延鋼板を得ることができる。本発明の高強度鋼板を自動車部品用途に使用すると、自動車の衝突安全性改善と燃費向上に大きく寄与することができる。
According to the present invention, it is possible to obtain a high-strength hot-rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and excellent strength, toughness and delayed fracture resistance after post-heating. When the high-strength steel sheet of the present invention is used for automobile parts, it can greatly contribute to improving the collision safety and fuel efficiency of automobiles.
According to the present invention, it is possible to obtain a high-strength hot-rolled steel sheet having a TS of 1180 MPa or more and less than 1600 MPa and excellent strength, toughness and delayed fracture resistance after post-heating. When the high-strength steel sheet of the present invention is used for automobile parts, it can greatly contribute to improving the collision safety and fuel efficiency of automobiles.
Claims (3)
- 質量％で、
Ｃ：０．０６～０．２３％、
Ｓｉ：０．１～３．０％、
Ｍｎ：１．５～３．５％、
Ｐ：０％超０．０５０％以下、
Ｓ：０％超０．００５０％以下、
Ａｌ：０％超１．５％以下、
Ｎ：０％超０．０１０％以下、
Ｏ：０％超０．００３％以下を含み、さらに、
ＴｉとＮｂを合計で０．０４０～０．２００％含み、
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、
鋼組織は、マルテンサイトおよび／または下部ベイナイトを主相とし、残留オーステナイトが体積率で３％未満であり、
Ｔｉ含有量とＮｂ含有量の合計に対する、固溶Ｔｉ量と固溶Ｎｂ量の合計の比である（固溶Ｔｉ量＋固溶Ｎｂ量）／（全Ｔｉ量＋全Ｎｂ量）が０．３００以上０．８００未満であり、
粒径１００ｎｍ以上の析出物として存在しているＴｉ量とＮｂ量が合計で０．０１０～０．０３０質量％であり、
表面から板厚中央方向に１００μｍまでの表層領域において、｛１１０｝＜１１１＞方位の極密度が１．８～５．０である、高強度熱延鋼板。 In mass percent,
C: 0.06 to 0.23%,
Si: 0.1 to 3.0%,
Mn: 1.5 to 3.5%,
P: more than 0% and not more than 0.050%;
S: more than 0% and 0.0050% or less;
Al: more than 0% and not more than 1.5%;
N: more than 0% and not more than 0.010%;
O: more than 0% and 0.003% or less;
Contains Ti and Nb in total 0.040 to 0.200%;
The balance is Fe and unavoidable impurities,
The steel structure has martensite and/or lower bainite as a main phase, and the volume fraction of retained austenite is less than 3%;
the ratio of the total amount of dissolved Ti and dissolved Nb to the total amount of Ti and Nb, that is, (amount of dissolved Ti+amount of dissolved Nb)/(total amount of Ti+total amount of Nb), is 0.300 or more and less than 0.800;
The total amount of Ti and Nb present as precipitates having a grain size of 100 nm or more is 0.010 to 0.030 mass %,
A high-strength hot-rolled steel sheet, in which the pole density of the {110}<111> orientation is 1.8 to 5.0 in a surface layer region extending from the surface to 100 μm in the sheet thickness center direction. - 前記成分組成が、さらに、質量％で、
Ｃｒ：０．００５～２．０％、
Ｎｉ：０．００５～２．０％、
Ｍｏ：０．００５～１．０％、
Ｖ：０．００５～０．５％、
Ｂ：０．０００２～０．００５０％、
Ｃａ：０．０００１～０．００５０％、
ＲＥＭ：０．０００１～０．００５０％
Ｃｕ：０．００５～０．５％、
Ｓｂ：０．００１０～０．１０％、および
Ｓｎ：０．００１０～０．１０％
のうちから選ばれる１種以上を含む、請求項１に記載の高強度熱延鋼板。 The composition further comprises, in mass%,
Cr: 0.005 to 2.0%,
Ni: 0.005 to 2.0%,
Mo: 0.005 to 1.0%,
V: 0.005 to 0.5%,
B: 0.0002 to 0.0050%,
Ca: 0.0001 to 0.0050%,
REM: 0.0001 to 0.0050%
Cu: 0.005 to 0.5%,
Sb: 0.0010 to 0.10%, and Sn: 0.0010 to 0.10%
The high strength hot rolled steel sheet according to claim 1, comprising at least one selected from the following: - 請求項１または２に記載の高強度熱延鋼板の製造方法であって、
前記成分組成を有するスラブを１１５０～１３００℃の温度域に加熱し、該温度域で０．２～３．５時間保持し、
次いで、熱間圧延を施すに際し、
１０８０℃以上の温度域での合計圧下率が８０～９０％、９００℃以下の温度域での合計圧下率が２０％以上、かつ、下記式で求められるＴ（℃）以下での１パスあたりの圧下率が２５％以下となる条件で圧延した後、１．０ｓ以上放冷し、
次いで、５５０℃までの温度域を５０℃／ｓ以上の平均冷却速度で冷却し、５５０℃に達してから急冷を開始するまでの時間を０．５～４．０ｓとし、
次いで、１００～２５０℃の巻取り温度までを２００℃／ｓ以上の冷却速度で急冷し、前記巻取り温度で巻取る、高強度熱延鋼板の製造方法。
Ｔ（℃）＝８００＋１０００［Ｔｉ］＋２５００［Ｎｂ］
ただし、［Ｔｉ］、［Ｎｂ］は、それぞれＴｉ、Ｎｂの含有量（質量％）であり、含有しない場合は０とする。
A method for producing a high strength hot rolled steel sheet according to claim 1 or 2,
A slab having the above-mentioned composition is heated to a temperature range of 1150 to 1300°C and held at that temperature range for 0.2 to 3.5 hours;
Next, when hot rolling is performed,
The total reduction in the temperature range of 1080°C or more is 80 to 90%, the total reduction in the temperature range of 900°C or less is 20% or more, and the reduction per pass at T (°C) or less calculated by the following formula is 25% or less. Then, the steel is allowed to cool for 1.0 s or more.
Next, the temperature is cooled at an average cooling rate of 50°C/s or more up to 550°C, and the time from reaching 550°C to starting quenching is set to 0.5 to 4.0 s.
Next, the steel sheet is quenched to a coiling temperature of 100 to 250° C. at a cooling rate of 200° C./s or more, and then coiled at the coiling temperature.
T(℃)=800+1000[Ti]+2500[Nb]
Here, [Ti] and [Nb] are the contents (mass%) of Ti and Nb, respectively, and are set to 0 when no Ti and Nb are contained.
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| WO2014132968A1 (en) * | 2013-02-26 | 2014-09-04 | 新日鐵住金株式会社 | HIGH-STRENGTH HOT-ROLLED STEEL SHEET HAVING MAXIMUM TENSILE STRENGTH OF 980 MPa OR ABOVE, AND HAVING EXCELLENT AND BAKING HARDENABILITY AND LOW-TEMPERATURE TOUGHNESS |
| WO2014185405A1 (en) * | 2013-05-14 | 2014-11-20 | 新日鐵住金株式会社 | Hot-rolled steel sheet and production method therefor |
| WO2014188966A1 (en) * | 2013-05-21 | 2014-11-27 | 新日鐵住金株式会社 | Hot-rolled steel sheet and method for manufacturing same |
| WO2018138887A1 (en) * | 2017-01-27 | 2018-08-02 | 新日鐵住金株式会社 | Steel sheet and plated steel sheet |
| WO2022042729A1 (en) * | 2020-08-31 | 2022-03-03 | 宝山钢铁股份有限公司 | 980 mpa-grade ultra-low-carbon martensite and retained austenite ultra-high hole expansion steel and manufacturing method therefor |
| WO2022042730A1 (en) * | 2020-08-31 | 2022-03-03 | 宝山钢铁股份有限公司 | High-strength low-carbon martensitic high hole expansion steel and manufacturing method therefor |
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| WO2014132968A1 (en) * | 2013-02-26 | 2014-09-04 | 新日鐵住金株式会社 | HIGH-STRENGTH HOT-ROLLED STEEL SHEET HAVING MAXIMUM TENSILE STRENGTH OF 980 MPa OR ABOVE, AND HAVING EXCELLENT AND BAKING HARDENABILITY AND LOW-TEMPERATURE TOUGHNESS |
| WO2014185405A1 (en) * | 2013-05-14 | 2014-11-20 | 新日鐵住金株式会社 | Hot-rolled steel sheet and production method therefor |
| WO2014188966A1 (en) * | 2013-05-21 | 2014-11-27 | 新日鐵住金株式会社 | Hot-rolled steel sheet and method for manufacturing same |
| WO2018138887A1 (en) * | 2017-01-27 | 2018-08-02 | 新日鐵住金株式会社 | Steel sheet and plated steel sheet |
| WO2022042729A1 (en) * | 2020-08-31 | 2022-03-03 | 宝山钢铁股份有限公司 | 980 mpa-grade ultra-low-carbon martensite and retained austenite ultra-high hole expansion steel and manufacturing method therefor |
| WO2022042730A1 (en) * | 2020-08-31 | 2022-03-03 | 宝山钢铁股份有限公司 | High-strength low-carbon martensitic high hole expansion steel and manufacturing method therefor |
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