JPWO2015118864A1 - Manufacturing method of high strength hot-rolled steel sheet - Google Patents
Manufacturing method of high strength hot-rolled steel sheet Download PDFInfo
- Publication number
- JPWO2015118864A1 JPWO2015118864A1 JP2015521173A JP2015521173A JPWO2015118864A1 JP WO2015118864 A1 JPWO2015118864 A1 JP WO2015118864A1 JP 2015521173 A JP2015521173 A JP 2015521173A JP 2015521173 A JP2015521173 A JP 2015521173A JP WO2015118864 A1 JPWO2015118864 A1 JP WO2015118864A1
- Authority
- JP
- Japan
- Prior art keywords
- less
- steel sheet
- hot
- rolled steel
- strength
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Granted
Links
- 229910000831 Steel Inorganic materials 0.000 title claims abstract description 209
- 239000010959 steel Substances 0.000 title claims abstract description 209
- 238000004519 manufacturing process Methods 0.000 title claims abstract description 13
- 229910000859 α-Fe Inorganic materials 0.000 claims abstract description 74
- 229910052719 titanium Inorganic materials 0.000 claims abstract description 22
- 239000002245 particle Substances 0.000 claims abstract description 19
- 239000013078 crystal Substances 0.000 claims abstract description 17
- 229910052748 manganese Inorganic materials 0.000 claims abstract description 15
- 229910052710 silicon Inorganic materials 0.000 claims abstract description 15
- 229910052758 niobium Inorganic materials 0.000 claims abstract description 14
- 239000012535 impurity Substances 0.000 claims abstract description 13
- 229910052799 carbon Inorganic materials 0.000 claims abstract description 11
- 229910052757 nitrogen Inorganic materials 0.000 claims abstract description 9
- 229910052782 aluminium Inorganic materials 0.000 claims abstract description 5
- 229910052698 phosphorus Inorganic materials 0.000 claims abstract description 3
- 238000005096 rolling process Methods 0.000 claims description 80
- 239000000463 material Substances 0.000 claims description 46
- 229910001566 austenite Inorganic materials 0.000 claims description 35
- 238000001816 cooling Methods 0.000 claims description 23
- 239000000203 mixture Substances 0.000 claims description 16
- 238000007747 plating Methods 0.000 claims description 11
- 229910052717 sulfur Inorganic materials 0.000 claims description 9
- 230000009467 reduction Effects 0.000 claims description 8
- 229910052718 tin Inorganic materials 0.000 claims description 5
- 229910052720 vanadium Inorganic materials 0.000 claims description 5
- 229910052750 molybdenum Inorganic materials 0.000 claims description 4
- 229910052725 zinc Inorganic materials 0.000 claims description 4
- 229910052787 antimony Inorganic materials 0.000 claims description 3
- 229910052785 arsenic Inorganic materials 0.000 claims description 3
- 229910052792 caesium Inorganic materials 0.000 claims description 3
- 229910052791 calcium Inorganic materials 0.000 claims description 3
- 229910052804 chromium Inorganic materials 0.000 claims description 3
- 229910052802 copper Inorganic materials 0.000 claims description 3
- 229910052745 lead Inorganic materials 0.000 claims description 3
- 229910052749 magnesium Inorganic materials 0.000 claims description 3
- 229910052759 nickel Inorganic materials 0.000 claims description 3
- 229910052761 rare earth metal Inorganic materials 0.000 claims description 3
- 229910052711 selenium Inorganic materials 0.000 claims description 3
- 229910052715 tantalum Inorganic materials 0.000 claims description 3
- 229910052721 tungsten Inorganic materials 0.000 claims description 3
- 229910052726 zirconium Inorganic materials 0.000 claims description 3
- 150000001247 metal acetylides Chemical class 0.000 abstract description 19
- 238000000034 method Methods 0.000 description 27
- 238000001953 recrystallisation Methods 0.000 description 24
- 238000005098 hot rolling Methods 0.000 description 18
- 238000004804 winding Methods 0.000 description 14
- 238000005204 segregation Methods 0.000 description 13
- 238000012360 testing method Methods 0.000 description 12
- 230000009466 transformation Effects 0.000 description 12
- 229910052751 metal Inorganic materials 0.000 description 10
- 239000002184 metal Substances 0.000 description 10
- 238000005728 strengthening Methods 0.000 description 9
- 230000000694 effects Effects 0.000 description 8
- 229910000734 martensite Inorganic materials 0.000 description 7
- 239000006104 solid solution Substances 0.000 description 7
- 239000002244 precipitate Substances 0.000 description 6
- 238000001556 precipitation Methods 0.000 description 6
- 230000015572 biosynthetic process Effects 0.000 description 5
- 238000004080 punching Methods 0.000 description 5
- 229910001563 bainite Inorganic materials 0.000 description 4
- 239000005539 carbonized material Substances 0.000 description 4
- 238000005266 casting Methods 0.000 description 4
- 230000007423 decrease Effects 0.000 description 4
- 238000005516 engineering process Methods 0.000 description 4
- 230000001376 precipitating effect Effects 0.000 description 4
- 230000008569 process Effects 0.000 description 4
- 238000012545 processing Methods 0.000 description 4
- 230000000717 retained effect Effects 0.000 description 4
- 238000009864 tensile test Methods 0.000 description 4
- 230000005540 biological transmission Effects 0.000 description 3
- 239000000470 constituent Substances 0.000 description 3
- 238000010438 heat treatment Methods 0.000 description 3
- 150000004767 nitrides Chemical class 0.000 description 3
- 150000003568 thioethers Chemical class 0.000 description 3
- 239000011701 zinc Substances 0.000 description 3
- 229910000838 Al alloy Inorganic materials 0.000 description 2
- 229910001335 Galvanized steel Inorganic materials 0.000 description 2
- 229910000746 Structural steel Inorganic materials 0.000 description 2
- 229910010413 TiO 2 Inorganic materials 0.000 description 2
- ATJFFYVFTNAWJD-UHFFFAOYSA-N Tin Chemical compound [Sn] ATJFFYVFTNAWJD-UHFFFAOYSA-N 0.000 description 2
- 229910045601 alloy Inorganic materials 0.000 description 2
- 239000000956 alloy Substances 0.000 description 2
- 239000002131 composite material Substances 0.000 description 2
- 238000009749 continuous casting Methods 0.000 description 2
- 238000005260 corrosion Methods 0.000 description 2
- 230000007797 corrosion Effects 0.000 description 2
- 230000001934 delay Effects 0.000 description 2
- 230000003111 delayed effect Effects 0.000 description 2
- 238000011161 development Methods 0.000 description 2
- 230000018109 developmental process Effects 0.000 description 2
- 238000007598 dipping method Methods 0.000 description 2
- 238000009713 electroplating Methods 0.000 description 2
- 239000008397 galvanized steel Substances 0.000 description 2
- 238000010191 image analysis Methods 0.000 description 2
- 230000001771 impaired effect Effects 0.000 description 2
- 230000006872 improvement Effects 0.000 description 2
- 238000002844 melting Methods 0.000 description 2
- 230000008018 melting Effects 0.000 description 2
- 229910001562 pearlite Inorganic materials 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 238000003825 pressing Methods 0.000 description 2
- 238000011160 research Methods 0.000 description 2
- 229910018137 Al-Zn Inorganic materials 0.000 description 1
- 229910018573 Al—Zn Inorganic materials 0.000 description 1
- HCHKCACWOHOZIP-UHFFFAOYSA-N Zinc Chemical compound [Zn] HCHKCACWOHOZIP-UHFFFAOYSA-N 0.000 description 1
- 230000002411 adverse Effects 0.000 description 1
- 238000005275 alloying Methods 0.000 description 1
- XAGFODPZIPBFFR-UHFFFAOYSA-N aluminium Chemical compound [Al] XAGFODPZIPBFFR-UHFFFAOYSA-N 0.000 description 1
- 238000000137 annealing Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 238000010276 construction Methods 0.000 description 1
- 238000005520 cutting process Methods 0.000 description 1
- 238000010586 diagram Methods 0.000 description 1
- 230000007613 environmental effect Effects 0.000 description 1
- 239000000446 fuel Substances 0.000 description 1
- 238000011835 investigation Methods 0.000 description 1
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
- 238000000465 moulding Methods 0.000 description 1
- 230000003287 optical effect Effects 0.000 description 1
- 238000005457 optimization Methods 0.000 description 1
- 230000008520 organization Effects 0.000 description 1
- 238000005498 polishing Methods 0.000 description 1
- 238000007670 refining Methods 0.000 description 1
- 238000010998 test method Methods 0.000 description 1
- 239000010409 thin film Substances 0.000 description 1
- 238000009849 vacuum degassing Methods 0.000 description 1
- 230000037303 wrinkles Effects 0.000 description 1
Images
Classifications
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- 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
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
- C21D8/0226—Hot rolling
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
- C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
- C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
- C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
-
- C—CHEMISTRY; METALLURGY
- C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
- C22C—ALLOYS
- C22C18/00—Alloys based on zinc
- C22C18/04—Alloys based on zinc with aluminium as the next major constituent
-
- 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
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/04—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the coating material
- C23C2/06—Zinc or cadmium or alloys based thereon
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/26—After-treatment
- C23C2/28—Thermal after-treatment, e.g. treatment in oil bath
-
- C—CHEMISTRY; METALLURGY
- C23—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; CHEMICAL SURFACE TREATMENT; DIFFUSION TREATMENT OF METALLIC MATERIAL; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL; INHIBITING CORROSION OF METALLIC MATERIAL OR INCRUSTATION IN GENERAL
- C23C—COATING METALLIC MATERIAL; COATING MATERIAL WITH METALLIC MATERIAL; SURFACE TREATMENT OF METALLIC MATERIAL BY DIFFUSION INTO THE SURFACE, BY CHEMICAL CONVERSION OR SUBSTITUTION; COATING BY VACUUM EVAPORATION, BY SPUTTERING, BY ION IMPLANTATION OR BY CHEMICAL VAPOUR DEPOSITION, IN GENERAL
- C23C2/00—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor
- C23C2/34—Hot-dipping or immersion processes for applying the coating material in the molten state without affecting the shape; Apparatus therefor characterised by the shape of the material to be treated
- C23C2/36—Elongated material
- C23C2/40—Plates; Strips
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/004—Dispersions; Precipitations
-
- C—CHEMISTRY; METALLURGY
- C21—METALLURGY OF IRON
- C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
- C21D2211/00—Microstructure comprising significant phases
- C21D2211/005—Ferrite
Landscapes
- Chemical & Material Sciences (AREA)
- Engineering & Computer Science (AREA)
- Organic Chemistry (AREA)
- Materials Engineering (AREA)
- Mechanical Engineering (AREA)
- Metallurgy (AREA)
- Thermal Sciences (AREA)
- Physics & Mathematics (AREA)
- Chemical Kinetics & Catalysis (AREA)
- Crystallography & Structural Chemistry (AREA)
- Oil, Petroleum & Natural Gas (AREA)
- Heat Treatment Of Sheet Steel (AREA)
- Coating With Molten Metal (AREA)
Abstract
本発明は、引張強さが590MPa以上であり、加工性に優れ、引張強さの異方性が小さい高加工性高強度熱延鋼板およびその製造方法を提供するものである。本発明に係る鋼板は、質量%で、C:0.010%超0.070%以下、Si:0.30%以下、Mn:0.30%超1.00%以下、P:0.030%以下、S:0.030%以下、Al:0.10%以下、N:0.0100%以下、Ti:0.050%以上0.120%以下を含有し、不純物元素であるNbおよびBをNb:0.005%以下、B:0.0005%以下に制限し、残部がFeおよび不可避的不純物からなる組成とし、フェライト相の面積率が95%以上であり、前記フェライト相の結晶粒の平均アスペクト比が3.0以下であり、前記フェライト相の結晶粒内にTiを含む炭化物が微細析出し、該炭化物の平均粒子径が10nm未満である組織とする。The present invention provides a high workability high strength hot-rolled steel sheet having a tensile strength of 590 MPa or more, excellent workability, and low tensile strength anisotropy, and a method for producing the same. The steel sheet according to the present invention is mass%, C: more than 0.010% and 0.070% or less, Si: 0.30% or less, Mn: more than 0.30% and 1.00% or less, P: 0.030. %: S: 0.030% or less, Al: 0.10% or less, N: 0.0100% or less, Ti: 0.050% or more and 0.120% or less, and Nb and B which are impurity elements Nb: 0.005% or less, B: 0.0005% or less, the balance is composed of Fe and inevitable impurities, the ferrite phase area ratio is 95% or more, the ferrite phase crystal grains The average aspect ratio is 3.0 or less, carbides containing Ti are finely precipitated in the ferrite phase crystal grains, and the carbide has an average particle diameter of less than 10 nm.
Description
本発明は、自動車をはじめとする輸送機械類の部品、建築用鋼材などの構造用鋼材に適した高強度熱延鋼板であって、引張強さ(TS):590MPa以上の高強度を有し、伸びフランジ性に優れると共に、引張強さの異方性が小さい高強度熱延鋼板およびその製造方法に関する。 The present invention is a high-strength hot-rolled steel sheet suitable for structural steel materials such as parts for transportation machinery including automobiles and building steel, and has a tensile strength (TS) of 590 MPa or more. The present invention relates to a high-strength hot-rolled steel sheet having excellent stretch flangeability and small anisotropy in tensile strength and a method for producing the same.
地球環境保全の観点から、CO2排出量を削減する目的で、自動車車体の強度を維持しつつその軽量化を図り自動車の燃費を改善することが、自動車業界において常に重要な課題とされている。自動車車体の強度を維持しつつ車体の軽量化を図るうえでは、自動車部品用素材となる鋼板の高強度化により、鋼板を薄肉化することが有効である。例えば、肉厚の鋼板を使用することの多い自動車の足回り部品用鋼板の薄肉化・高強度化により、自動車車体の大幅な軽量化が期待できる。From the viewpoint of global environmental conservation, for the purpose of reducing CO 2 emissions, maintaining the strength of the car body while reducing the weight and improving the fuel efficiency of the car has always been an important issue in the automobile industry. . In order to reduce the weight of the vehicle body while maintaining the strength of the automobile body, it is effective to reduce the thickness of the steel sheet by increasing the strength of the steel sheet used as a material for automobile parts. For example, by reducing the thickness and strength of steel plates for undercarriage parts of automobiles, which often use thick steel plates, it is possible to expect a significant reduction in the weight of automobile bodies.
一方、鋼板を素材とする自動車部品の多くは、プレス加工やバーリング加工等によって成形されるため、自動車部品用鋼板には優れた成形性が要求される。例えば、自動車の足回り用部品では、その部品形状からバーリング加工を施すことが多く、優れた伸びフランジ性を有する鋼板が必要とされる。 On the other hand, since many automobile parts made of steel plates are formed by press working or burring, etc., excellent formability is required for steel sheets for automobile parts. For example, parts for automobile undercarriages are often subjected to burring due to the shape of the parts, and steel plates having excellent stretch flangeability are required.
また、引張強さの異方性が大きい鋼板にプレス加工やバーリング加工を施す場合には、この異方性に起因するスプリングバック量の変動が大きくなってしまうため、鋼板から部品を打ち抜く際の方向が限定される。これに対し、引張強さの異方性が小さい鋼板では、異方性に起因するスプリングバック量の変動が小さく、鋼板から部品を打ち抜く際の方向が限定されないため、大幅な歩留まりの向上が期待できる。それゆえ、鋼板の引張強さの異方性を小さくすることも、産業上極めて重要である。特に、自動車足回り用部品のように、複雑な形状の部品を打ち抜く際、歩留まりの向上は顕著となる。 In addition, when pressing or burring a steel sheet with a large tensile strength anisotropy, fluctuations in the amount of springback due to this anisotropy will increase. The direction is limited. In contrast, steel sheets with low anisotropy in tensile strength have little variation in the amount of springback due to anisotropy, and the direction in which parts are punched from the steel sheet is not limited, so a significant improvement in yield is expected. it can. Therefore, it is extremely important in the industry to reduce the anisotropy of the tensile strength of the steel sheet. In particular, when punching out parts having complicated shapes such as parts for automobile undercarriage, the yield is significantly improved.
以上の理由により、特に自動車足回り部品に適用される鋼板としては、伸びフランジ性に優れると共に、引張強さの異方性が小さい高強度鋼板が望まれている。しかしながら、一般的に鉄鋼材料は高強度化に伴い延性や伸びフランジ性が低下し、更に機械的特性の異方性が大きくなる傾向にある。このため、高強度鋼板を自動車の足回り部品に適用するにあたっては、高強度であることに加え、伸びフランジ性に優れ、更に機械的特性の異方性が小さい鋼板が必要とされる。 For the reasons described above, a high-strength steel sheet that is excellent in stretch flangeability and has a small tensile strength anisotropy is particularly desired as a steel sheet applied to automobile undercarriage parts. However, in general, steel materials tend to have lower ductility and stretch flangeability with increasing strength, and further increase the anisotropy of mechanical properties. For this reason, when a high-strength steel plate is applied to an automobile undercarriage part, a steel plate that is excellent in stretch flangeability and has low mechanical property anisotropy is required in addition to high strength.
強度と加工性を兼ね備えた高強度熱延鋼板に関しては、数多くの研究開発が為され、各種技術が提案されている。中でも、金属組織を実質的にフェライト単相とし、フェライト相の粒内に微細炭化物を析出させることにより鋼板を高強度化する技術は、高強度でありながら優れた伸びフランジ性を両立させるために非常に有用な技術であることが知られている。 Many researches and developments have been made on high-strength hot-rolled steel sheets that have both strength and workability, and various technologies have been proposed. Above all, the technology to increase the strength of the steel sheet by making the metal structure substantially ferrite single phase and precipitating fine carbides in the ferrite phase grains is to achieve both excellent strength and excellent flange flangeability. It is known to be a very useful technique.
例えば、特許文献1には、熱延鋼板に関し、鋼板組成をwt%でC:0.01〜0.10%、Si:1.5%以下、Mn:1.0%超〜2.5%、P:0.15%以下、S:0.008%以下、Al:0.01〜0.08%、Ti、Nbの1種または2種の合計:0.10〜0.60%を含み、Feを主成分とする組成とし、鋼板組織をフェライト量が面積率で95%以上、かつフェライトの平均結晶粒径が2.0〜10.0μmであり、マルテンサイトおよび残留オーステナイトを実質的に含まない組織とする技術が提案されている。そして、特許文献1には、Ti、Nbの1種または2種の合計含有量を0.10〜0.60%とすることで熱延過程での再結晶が抑制され、更にMn含有量を1.0%超〜2.5%とすることでフェライト粒の粗大化が抑制される結果、微細フェライト粒が得られ、伸びフランジ性を損なうことなく引張強さ490MPa以上の熱延鋼板が得られると記載されている。 For example, Patent Document 1 relates to a hot-rolled steel sheet, in which the steel sheet composition is wt%, C: 0.01 to 0.10%, Si: 1.5% or less, Mn: more than 1.0% to 2.5% , P: 0.15% or less, S: 0.008% or less, Al: 0.01 to 0.08%, total of one or two of Ti and Nb: 0.10 to 0.60% The composition of which the main component is Fe, the steel sheet structure has a ferrite content of 95% or more by area ratio, and the average crystal grain size of ferrite is 2.0 to 10.0 μm, and substantially contains martensite and retained austenite. A technology that does not include the organization has been proposed. And in patent document 1, recrystallization in a hot rolling process is suppressed by making 1 type or 2 types of total content of Ti and Nb into 0.10 to 0.60%, and also Mn content is made into As a result of suppressing the coarsening of the ferrite grains by setting it to more than 1.0% to 2.5%, fine ferrite grains are obtained, and a hot rolled steel sheet having a tensile strength of 490 MPa or more is obtained without impairing stretch flangeability. It is stated that
また、特許文献2には、熱延鋼板に関し、鋼板組成を質量%でC:0.01〜0.1%、S:0.03%以下、N:0.005%以下、Ti:0.05〜0.5%を含み、さらに(Ti−48/12C−48/14N−48/32S)≧0%を満たす範囲でTiを含有し、残部がFeおよび不可避的不純物かなる組成とし、鋼中の粒子で5nm以上のTiを含む析出物の平均サイズを101〜103nm、最小間隔を101nm超104nm以下とする技術が提案されている。そして、特許文献2には、5nm以上のTiを含む析出物の平均サイズと最小間隔を上記の如く規定することにより、バーリング加工性と疲労特性に優れた引張強度640MPa以上の熱延鋼板が得られると記載されている。In addition,
一方、異方性の小さな高強度熱延鋼板を製造する技術に関しても多くの研究開発が為されている。例えば、特許文献3には、熱延鋼板に関し、フェライトまたはベイナイトを体積分率最大の相とし、体積分率で1%以上25%以下のマルテンサイトや残留オーステナイトを任意に含み、少なくとも1/2板厚における板面の{100}<011>〜{223}<110>方位群のX線ランダム強度比の平均値を2.5以上で、かつ、{554}<225>、{111}<112>および{111}<110>の3つの結晶方位のX線ランダム強度比の平均値を3.5以下とし、更に、圧延方向のr値および圧延方向と直角方向のr値のうち少なくとも1つを0.7以下とし、均一伸びの異方性ΔuElを4%以下で、かつ、局部伸びの異方性ΔLEl以下とする技術が提案されている。 On the other hand, many researches and developments have been made on techniques for producing high-strength hot-rolled steel sheets with small anisotropy. For example, Patent Document 3 relates to a hot-rolled steel sheet, with ferrite or bainite being the phase with the largest volume fraction, optionally containing martensite or residual austenite with a volume fraction of 1% to 25%, and at least 1/2. The average value of the X-ray random intensity ratio of the {100} <011> to {223} <110> orientation groups on the plate thickness in the plate thickness is 2.5 or more, and {554} <225>, {111} < The average value of the X-ray random intensity ratios of the three crystal orientations 112> and {111} <110> is 3.5 or less, and at least one of the r value in the rolling direction and the r value in the direction perpendicular to the rolling direction. A technique has been proposed in which one is 0.7 or less, the uniform elongation anisotropy ΔuE1 is 4% or less, and the local elongation anisotropy ΔLE1 is less than or equal to.
そして、特許文献3には、上記の如く特定の集合組織とし、且つΔuElとΔLElを規定することにより、形状凍結性に優れた異方性の小さい高加工性高強度熱延鋼板が得られると記載されている。 In Patent Document 3, a specific texture as described above, and by defining ΔuEl and ΔLEl, a highly workable and high strength hot rolled steel sheet having low shape anisotropy and low anisotropy can be obtained. Have been described.
しかしながら、特許文献1に提案された技術では、鋼板に1.0%超のMnを添加する必要があるため、鋼板中にMnの偏析が不可避的に発生する。このようにMnの偏析が発生した鋼板では、Mn偏析部の強度が非偏析部の強度に比べて著しく上昇するため、加工時にMn偏析部を起点とした亀裂が発生し、鋼板の伸びフランジ性の安定性に問題がある。また、特許文献1に提案された技術では、Ti、Nbの1種または2種を合計で0.10〜0.60%含有させて熱延過程での再結晶を抑制することで、フェライト粒の細粒化を図り、鋼板の強度と伸びフランジ性を確保している。このように熱延過程での再結晶を抑制すると、仕上げ圧延後の再結晶が抑制される結果、変態後のフェライト粒形状が扁平となり、引張強さの異方性が大きくなるという問題がある。そのため、特許文献1に提案された技術では、引張強さの異方性を抑制しつつ、鋼板の強度と伸びフランジ性を確保することができない。 However, in the technique proposed in Patent Document 1, it is necessary to add more than 1.0% of Mn to the steel sheet, so Mn segregation inevitably occurs in the steel sheet. Thus, in the steel sheet in which Mn segregation occurs, the strength of the Mn segregation part is significantly higher than the strength of the non-segregation part. There is a problem with stability. Further, in the technique proposed in Patent Document 1, ferrite grains are contained by containing 0.10 to 0.60% of one or two of Ti and Nb in total to suppress recrystallization in the hot rolling process. The strength of the steel sheet and stretch flangeability are ensured. In this way, if recrystallization in the hot rolling process is suppressed, recrystallization after finish rolling is suppressed, resulting in a problem that the ferrite grain shape after transformation becomes flat and the anisotropy of tensile strength increases. . Therefore, the technique proposed in Patent Document 1 cannot secure the strength and stretch flangeability of the steel sheet while suppressing the anisotropy of the tensile strength.
また、特許文献2に提案された技術では、(Ti−48/12C−48/14N−48/32S)≧0%を満たす範囲でTiを含有させることにより炭化物を形成し、固溶Cの低減化を図っている。このようにCに対して過剰のTiを含有させた場合、熱延鋼板製造工程の巻取り温度近傍において固溶Ti量が著しく増大し、巻き取り後のTi炭化物の粗大化が促進される。したがって、特許文献2に提案された技術では、製造条件(主として巻取り温度)の変動に対する鋼板強度の安定性が、著しく低下するという問題がある。
Further, in the technique proposed in
特許文献3に提案された技術では、鋼板の長手方向、幅方向に安定して特定の集合組織を得ることが困難である。また、特許文献3に提案された技術では、鋼板組織としてマルテンサイトおよび残留オーステナイトを積極的に含有させるため、熱延条件に対する強度安定性が著しく低くなるという問題がある。更に、特許文献3に提案された技術により得られる熱延鋼板は、フェライトまたはベイナイトに加えてマルテンサイトおよび残留オーステナイトを積極的に含有させた複合組織を有するため、マルテンサイトなどの硬質相とフェライト相との界面に硬度差に起因する亀裂が発生し易く、伸びフランジ性に優れる鋼板とはいい難い。 With the technique proposed in Patent Document 3, it is difficult to obtain a specific texture stably in the longitudinal direction and the width direction of the steel sheet. Moreover, in the technique proposed in Patent Document 3, since martensite and retained austenite are positively contained as a steel sheet structure, there is a problem that strength stability against hot rolling conditions is remarkably lowered. Furthermore, since the hot-rolled steel sheet obtained by the technique proposed in Patent Document 3 has a composite structure in which martensite and retained austenite are positively contained in addition to ferrite or bainite, a hard phase such as martensite and ferrite Cracks due to hardness differences are likely to occur at the interface with the phase, and it is difficult to say a steel plate with excellent stretch flangeability.
以上のように、従来技術では、強度と伸びフランジ性を確保しつつ、引張強さの異方性の小さい高強度熱延鋼板を得ることが極めて困難である。 As described above, according to the prior art, it is extremely difficult to obtain a high-strength hot-rolled steel sheet having small tensile strength anisotropy while ensuring strength and stretch flangeability.
本発明は、かかる従来技術の問題点を有利に解決し、自動車部品用の素材として好適な、引張強さ(TS):590MPa以上の高強度と優れた伸びフランジ性を有し、引張強さの異方性の小さい高強度熱延鋼板およびその製造方法を提供することを目的とする。なお、本発明において単に引張強さと表記した場合、C方向(圧延直角方向)の引張強さを表す。 The present invention advantageously solves the problems of the prior art and is suitable as a material for automobile parts. Tensile strength (TS): High strength of 590 MPa or more, excellent stretch flangeability, and tensile strength An object of the present invention is to provide a high-strength hot-rolled steel sheet having a small anisotropy and a method for producing the same. In addition, when it only describes with tensile strength in this invention, the tensile strength of C direction (rolling perpendicular direction) is represented.
上記課題を解決すべく、本発明者らは、加工性が良好なフェライト単相組織を有する熱延鋼板に着目し、優れた加工性を保ちつつ、熱延鋼板を高強度化し、更に引張強さの異方性を小さくする手法について鋭意検討した。 In order to solve the above-mentioned problems, the present inventors focused on hot-rolled steel sheets having a ferrite single-phase structure with good workability, and increased the strength of the hot-rolled steel sheets while maintaining excellent workability. The method for reducing the anisotropy of the thickness was studied earnestly.
フェライト単相組織の熱延鋼板を高強度化する手法としては、SiやMnによる固溶強化と、TiやNbによる析出強化が挙げられる。しかしながら、従来、鋼板の高強度化に極めて有用とされ、固溶強化元素として高強度鋼板に積極添加されてきたSiやMnは、板厚中央部に偏析した場合に伸びフランジ性の安定性を低下させる原因となる。また、TiやNbの析出により高強度化した鋼板では、熱延鋼板の製造工程において、TiやNbが仕上げ圧延後の再結晶を阻害するためにフェライト粒が扁平となり、機械的特性の異方性が大きくなり易い。 As a method for increasing the strength of a hot-rolled steel sheet having a ferrite single phase structure, solid solution strengthening by Si or Mn and precipitation strengthening by Ti or Nb can be mentioned. However, Si and Mn, which have been extremely useful for increasing the strength of steel sheets and have been positively added to high-strength steel sheets as solid solution strengthening elements, have improved stretch flangeability when segregated at the center of the plate thickness. It causes a decrease. Also, in steel sheets with high strength due to precipitation of Ti and Nb, in the manufacturing process of hot-rolled steel sheets, the ferrite grains become flat because Ti and Nb inhibit recrystallization after finish rolling, and the mechanical properties are anisotropic. Tend to be large.
このような問題に対し、本発明者らは、MnおよびSiを添加した熱延鋼板について組織観察し、MnおよびSiの含有量と伸びフランジ性との関係について、綿密な調査を行った。その結果、MnおよびSiが合計で1.5%を超えて含まれる場合には、板厚中央部に不可避的に偏析が存在し、偏析に起因する組織の硬さや形状の変化が鋼板の伸びフランジ性に悪影響を及ぼすという知見を得た。そして、MnおよびSiの含有量を所定量以下、具体的にはSi含有量を0.30%以下、Mn含有量を1.00%以下とすることで、上記偏析組織の影響を抑制できるという知見を得た。 In order to solve such a problem, the present inventors have observed the structure of hot rolled steel sheets to which Mn and Si are added, and have conducted a thorough investigation on the relationship between the contents of Mn and Si and stretch flangeability. As a result, when Mn and Si are included in excess of 1.5% in total, segregation inevitably exists in the central portion of the plate thickness, and changes in the hardness and shape of the structure due to segregation cause elongation of the steel sheet. The knowledge that it has a bad influence on flangeability was obtained. And, it is said that the influence of the segregated structure can be suppressed by setting the contents of Mn and Si to a predetermined amount or less, specifically, the Si content to 0.30% or less and the Mn content to 1.00% or less. Obtained knowledge.
これらの知見に基づき、本発明者らは、SiおよびMnによる固溶強化で鋼板を高強度化する手法により、優れた伸びフランジ性と高強度を兼備しながら鋼板の機械的特性の異方性を小さくすることは極めて困難であると判断した。そして、フェライト単相組織の熱延鋼板を高強度化する手法として、SiおよびMnの添加量を低減しつつ、TiおよびNbの炭化物による析出強化を活用することを試みた。 Based on these findings, the present inventors have developed an anisotropy of the mechanical properties of the steel sheet while combining excellent stretch flangeability and high strength by means of increasing the strength of the steel sheet by solid solution strengthening with Si and Mn. It was judged that it was extremely difficult to reduce the size. As a technique for increasing the strength of a hot rolled steel sheet having a ferrite single phase structure, an attempt was made to utilize precipitation strengthening by carbides of Ti and Nb while reducing the amount of Si and Mn added.
鋼板にTiおよびNbの微細炭化物を析出させることにより、鋼板強度の大幅な向上が期待できる。しかしながら、炭化物形成元素であるTiやNbの添加は、熱延鋼板の製造工程において、仕上げ圧延後のオーステナイトの再結晶を遅延させる要因となる。オーステナイトの再結晶が促進されない場合には、γ→α変態前のオーステナイトが圧延方向に伸長した組織となる。その結果、変態後のフェライトも圧延方向に伸長した組織となり、圧延方向と圧延直角方向の機械的特性に顕著な差異が生じる。 By precipitating fine carbides of Ti and Nb on the steel plate, significant improvement in steel plate strength can be expected. However, the addition of Ti and Nb, which are carbide forming elements, is a factor that delays the recrystallization of austenite after finish rolling in the manufacturing process of hot-rolled steel sheets. When the recrystallization of austenite is not promoted, the austenite before the γ → α transformation has a structure elongated in the rolling direction. As a result, the ferrite after transformation also has a structure elongated in the rolling direction, and there is a marked difference in mechanical properties between the rolling direction and the direction perpendicular to the rolling direction.
そこで、本発明者らは更に検討を進め、所望の鋼板強度を得るに十分な量の炭化物形成元素を添加した場合においても、熱間圧延後のフェライト粒を等軸とする手段を模索した。その結果、炭化物形成元素としてNbを用いた場合ではフェライト粒を等軸とすることは困難であるが、炭化物形成元素としてTiを用いるとともに、SiおよびMnの添加量を所定量以下に低減し、且つ熱間圧延における仕上げ圧延の圧下率、圧延温度および圧延後保持時間を制御することにより、フェライト粒が等軸または等軸に近い形状となり、所望の強度と優れた伸びフランジ性を維持しつつ、引張強さの異方性が抑制可能であるという知見を得た。 Accordingly, the present inventors have further studied and sought a means for making the ferrite grains after hot rolling equiaxed even when a sufficient amount of carbide forming element is added to obtain a desired steel plate strength. As a result, when Nb is used as the carbide forming element, it is difficult to make the ferrite grains equiaxed, but while using Ti as the carbide forming element, the addition amount of Si and Mn is reduced to a predetermined amount or less, In addition, by controlling the reduction ratio of finish rolling in hot rolling, the rolling temperature and the holding time after rolling, the ferrite grains become equiaxed or nearly equiaxed, while maintaining desired strength and excellent stretch flangeability. In addition, the inventors have found that the anisotropy of tensile strength can be suppressed.
本発明は上記の知見に基づき完成されたものであり、その要旨は次のとおりである。
[1]質量%で、C:0.010%超0.070%以下、Si:0.30%以下、Mn:0.30%超1.00%以下、P:0.030%以下、S:0.030%以下、Al:0.10%以下、N:0.0100%以下、Ti:0.050%以上0.120%以下を含有し、不純物元素であるNbおよびBを、質量%で、Nb:0.005%以下、B:0.0005%以下に制限し、残部がFeおよび不可避的不純物からなる組成を有し、フェライト相の面積率が95%以上であり、前記フェライト相の結晶粒の平均アスペクト比が3.0以下であり、前記フェライト相の結晶粒内にTiを含む炭化物が微細析出し、該炭化物の平均粒子径が10nm未満である組織を有し、引張強さが590MPa以上であることを特徴とする高強度熱延鋼板。
[2]前記[1]において、前記組成が、以下の(1)式を満足することを特徴とする高強度熱延鋼板。
((Ti−(48/14)×N−(48/32)×S)/48)/(C/12)<1.0 ・・・(1)
((1)式中のC、S、N、Ti:各元素の含有量(質量%))
[3]前記[1]または[2]において、前記組成に加えて更に、質量%で、REM、Zr、V、As、Cu、Ni、Sn、Pb、Ta、W、Mo、Cr、Sb、Mg、Ca、Co、Se、Zn、Csのうちから選ばれた1種以上を合計で1.0%以下含有することを特徴とする高強度熱延鋼板。
[4]前記[1]ないし[3]のいずれかにおいて、鋼板表面にめっき層を有することを特徴とする高強度熱延鋼板。
[5]前記[1]ないし[3]のいずれかに記載の組成からなる鋼素材を、オーステナイト単相域に加熱し、熱間圧延を施した後、冷却し、巻き取り、熱延鋼板とするにあたり、前記熱間圧延の仕上げ圧延において、900℃以上1100℃以下の温度域で圧下率15%以上の圧延を3回以上行い、仕上げ圧延温度を860℃以上とし、前記仕上げ圧延終了後、850℃以上の温度域で0.3s以上保持した後前記冷却を開始し、前記冷却の850℃から750℃までの平均冷却速度を30℃/s以上とし、前記巻き取りの巻取り温度を580℃以上750℃以下とすることを特徴とする高強度熱延鋼板の製造方法。The present invention has been completed based on the above findings, and the gist thereof is as follows.
[1] By mass%, C: more than 0.010% and 0.070% or less, Si: 0.30% or less, Mn: more than 0.30% and 1.00% or less, P: 0.030% or less, S : 0.030% or less, Al: 0.10% or less, N: 0.0100% or less, Ti: 0.050% or more and 0.120% or less, and Nb and B which are impurity elements are contained in mass%. Nb: 0.005% or less, B: 0.0005% or less, the balance is composed of Fe and inevitable impurities, the area ratio of the ferrite phase is 95% or more, the ferrite phase The crystal grains have an average aspect ratio of 3.0 or less, carbides containing Ti are finely precipitated in the ferrite phase grains, and the carbide has an average particle diameter of less than 10 nm. High-strength hot-rolled steel characterized by having a thickness of 590 MPa or more .
[2] A high-strength hot-rolled steel sheet according to [1], wherein the composition satisfies the following formula (1):
((Ti− (48/14) × N− (48/32) × S) / 48) / (C / 12) <1.0 (1)
(C, S, N, Ti in formula (1): content of each element (mass%))
[3] In the above [1] or [2], in addition to the composition, in addition to REM, Zr, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, A high-strength hot-rolled steel sheet containing 1.0% or less in total of one or more selected from Mg, Ca, Co, Se, Zn, and Cs.
[4] The high-strength hot-rolled steel sheet according to any one of [1] to [3], wherein the steel sheet surface has a plating layer.
[5] A steel material having the composition according to any one of [1] to [3] is heated to an austenite single-phase region, hot-rolled, cooled, wound, hot-rolled steel sheet and In doing so, in the finish rolling of the hot rolling, rolling at a rolling reduction of 15% or more in a temperature range of 900 ° C. or more and 1100 ° C. or less is performed three or more times, the finish rolling temperature is set to 860 ° C. or more, and after the finish rolling is finished, The cooling is started after being held for 0.3 s or more in a temperature range of 850 ° C. or more, the average cooling rate from 850 ° C. to 750 ° C. of the cooling is 30 ° C./s or more, and the winding temperature of the winding is 580 A method for producing a high-strength hot-rolled steel sheet, characterized by having a temperature of from ℃ to 750 ℃.
本発明によると、自動車をはじめとする輸送機械類の部品、建築用鋼材などの構造用鋼材に適した、引張強さ(TS):590MPa以上の高強度と優れた加工性を有し、且つ引張強さの異方性が小さい高強度熱延鋼板が得られる。また、上記の如く優れた鋼板特性が得られることから、本発明は、高強度熱延鋼板の更なる用途展開を可能とし、産業上格段の効果を奏する。 According to the present invention, the tensile strength (TS): suitable for structural steel materials such as parts for transportation machinery including automobiles and construction steel materials, has a high strength of 590 MPa or more and excellent workability, and A high-strength hot-rolled steel sheet with a small tensile strength anisotropy is obtained. Moreover, since the outstanding steel plate characteristic is acquired as mentioned above, this invention enables the further use expansion | deployment of a high intensity | strength hot-rolled steel plate, and has an industrial remarkable effect.
以下、本発明について具体的に説明する。 Hereinafter, the present invention will be specifically described.
まず、本発明熱延鋼板の成分組成の限定理由について説明する。なお、以下の成分組成を表す%は、特に断らない限り質量%(mass%)を意味するものとする。 First, the reasons for limiting the component composition of the hot-rolled steel sheet of the present invention will be described. In addition,% which represents the following component composition shall mean the mass% (mass%) unless there is particular notice.
C:0.010%超0.070%以下
Cは、鋼板中でTiを含む炭化物を形成し、熱延鋼板を高強度化するうえで必須の元素である。C含有量が0.010%以下であると、鋼板の引張強さを590MPa以上とするに十分な炭化物析出量を得ることができず、590MPa以上の引張強さが得られなくなる。一方、C含有量が0.070%を超えると、パーライトが生成し易くなり、鋼板の伸びフランジ性が低下する。したがって、C含有量は0.010%超0.070%以下とする。好ましくは、0.020%以上0.060%以下である。C: more than 0.010% and 0.070% or less C is an essential element for forming carbide containing Ti in the steel sheet and increasing the strength of the hot-rolled steel sheet. When the C content is 0.010% or less, a carbide precipitation amount sufficient to make the tensile strength of the steel sheet 590 MPa or more cannot be obtained, and a tensile strength of 590 MPa or more cannot be obtained. On the other hand, when the C content exceeds 0.070%, pearlite is easily generated, and the stretch flangeability of the steel sheet is deteriorated. Therefore, the C content is more than 0.010% and 0.070% or less. Preferably, it is 0.020% or more and 0.060% or less.
Si:0.30%以下
Siは、延性(伸び)低下をもたらすことなく鋼板強度を向上させる有効な元素として、通常、高強度鋼板に積極的に含有されている。しかしながら、Siは、本発明の熱延鋼板において回避すべき板厚中央部のMn偏析を助長するとともに、Si自身も偏析する元素である。したがって、本発明では、上記Mn偏析を抑制し、またSiの偏析を抑制する目的で、Si含有量を0.30%以下に限定する。好ましくは0.10%以下であり、より好ましくは0.05%以下である。Si: 0.30% or less Si is normally positively contained in a high-strength steel sheet as an effective element for improving the steel sheet strength without reducing ductility (elongation). However, Si is an element that promotes Mn segregation at the center of the plate thickness to be avoided in the hot-rolled steel sheet of the present invention and also segregates Si itself. Therefore, in the present invention, the Si content is limited to 0.30% or less for the purpose of suppressing the Mn segregation and suppressing the segregation of Si. Preferably it is 0.10% or less, More preferably, it is 0.05% or less.
Mn:0.30%超1.00%以下
Mnは、固溶強化元素であり、Siと同様、通常の高強度鋼板には積極的に含有されている。しかしながら、鋼板にMnを積極的に含有させると、板厚中央部のMn偏析は避けられず、鋼板の伸びフランジ性が劣化する原因となる。したがって、本発明では、上記Mn偏析を抑制することを目的として、Mn含有量を1.00%以下に限定する。好ましくは0.90%以下である。Mn: more than 0.30% and 1.00% or less Mn is a solid solution strengthening element, and is actively contained in a normal high-strength steel sheet like Si. However, if Mn is positively contained in the steel sheet, Mn segregation at the center of the plate thickness is unavoidable, which causes the stretch flangeability of the steel sheet to deteriorate. Therefore, in the present invention, the Mn content is limited to 1.00% or less for the purpose of suppressing the Mn segregation. Preferably it is 0.90% or less.
一方、Mn含有量が0.30%以下になると、γ→α変態点が上昇するため、Tiを含む炭化物の微細化が困難となる。Tiを含む炭化物は熱延鋼板製造工程における仕上げ圧延終了後の冷却または巻き取り過程でγ→α変態と同時に、もしくはフェライト中に時効析出する。ここで、γ→α変態点が高温になると、Tiを含む炭化物が高温域で析出することになるため、炭化物が粗大化してしまう。そして、Tiを含む炭化物が粗大化する結果、所望の鋼板強度が得られなくなる。したがって、Mn含有量は0.30%超とする。好ましくは0.35%超である。 On the other hand, when the Mn content is 0.30% or less, the γ → α transformation point increases, and it is difficult to refine the carbide containing Ti. The carbide containing Ti is aged at the same time as the γ → α transformation or in the ferrite during the cooling or winding process after finish rolling in the hot rolled steel sheet manufacturing process. Here, when the [gamma]-> [alpha] transformation point becomes high temperature, the carbide containing Ti is precipitated in a high temperature region, so that the carbide becomes coarse. And as a result of the carbide | carbonized_material containing Ti coarsening, desired steel plate intensity | strength will no longer be obtained. Therefore, the Mn content is more than 0.30%. Preferably it is more than 0.35%.
P:0.030%以下
Pは、粒界に偏析して鋼板の伸びを低下させ、加工時に割れを誘発する有害な元素である。したがって、P含有量は0.03%以下とする。P: 0.030% or less P is a harmful element that segregates at grain boundaries to reduce the elongation of the steel sheet and induces cracks during processing. Therefore, the P content is 0.03% or less.
S:0.030%以下
Sは、鋼中にMnSやTiSとして存在する。MnSやTiSは、熱延鋼板の打抜き加工時にボイドの発生を助長し、更には、加工中にもボイドの発生の起点となるため、鋼板の伸びフランジ性を低下させる。したがって、本発明では、S含有量を極力低減することが好ましく、0.030%以下とする。好ましくは0.010%以下である。S: 0.030% or less S is present in the steel as MnS or TiS. MnS and TiS promote the generation of voids during the punching process of hot-rolled steel sheets, and further serve as a starting point for the generation of voids during the processing, thereby reducing the stretch flangeability of the steel sheet. Therefore, in the present invention, it is preferable to reduce the S content as much as possible, and set it to 0.030% or less. Preferably it is 0.010% or less.
Al:0.10%以下
Alは、脱酸剤として作用する元素である。このような効果を得るためには、Alを0.01%以上含有することが望ましい。しかし、Al含有量が0.10%を超えると、Alが鋼板中にAl酸化物として残存し、該Al酸化物が凝集して粗大化し易くなり、鋼板の伸びフランジ性を劣化させる原因となる。したがって、Al含有量は0.10%以下とする。好ましくは0.05%以下である。Al: 0.10% or less Al is an element that acts as a deoxidizer. In order to obtain such an effect, it is desirable to contain 0.01% or more of Al. However, when the Al content exceeds 0.10%, Al remains as an Al oxide in the steel sheet, and the Al oxide tends to agglomerate and become coarse, which causes the stretch flangeability of the steel sheet to deteriorate. . Therefore, the Al content is set to 0.10% or less. Preferably it is 0.05% or less.
N:0.0100%以下
Nは、鋼中にTiNとして存在するため、N含有量が多くなると、炭化物の形成に寄与するTi量がNの存在により低下し、所望の鋼板強度が得られなくなる。加えて、TiNは熱延鋼板の打抜き加工時にボイドの発生を助長し、更には、加工中にもボイドの発生の起点となるため、鋼板の伸びフランジ性を低下させる。以上の理由により、本発明ではN含有量を極力低減することが好ましく、0.0100%以下とする。好ましくは0.0060%以下である。N: 0.0100% or less Since N is present as TiN in the steel, if the N content increases, the amount of Ti that contributes to the formation of carbides decreases due to the presence of N, and the desired steel plate strength cannot be obtained. . In addition, TiN promotes the generation of voids during the punching process of hot-rolled steel sheets, and further, since it becomes a starting point for the generation of voids during processing, it reduces the stretch flangeability of the steel sheet. For the above reasons, in the present invention, it is preferable to reduce the N content as much as possible, and it is set to 0.0100% or less. Preferably it is 0.0060% or less.
Ti:0.050%以上0.120%以下
Tiは、Tiを含む炭化物を形成して鋼板の高強度化を図るうえで必要不可欠な元素である。Ti含有量が0.050%未満では、所望の熱延鋼板強度(引張強さ:590MPa以上)を得ることが困難となる。一方、Ti含有量が過剰になるとTiを含む炭化物が粗大化する傾向が見られ、所望の熱延鋼板強度(引張強さ:590MPa以上)を得ることが困難となる。また、Tiは、オーステナイトの再結晶を阻害する元素である。そのため、Tiが0.120%を超えて多量に含有する場合には、仕上げ圧延後の再結晶が遅延し、変態後のフェライト組織が圧延方向に伸長し、鋼板の機械的特性の異方性が大きくなる。したがって、Ti含有量は0.050%以上0.120%以下とする。好ましくは0.060%以上0.110%以下、より好ましくは0.060%以上0.100%以下である。Ti: 0.050% or more and 0.120% or less Ti is an indispensable element for forming a carbide containing Ti and increasing the strength of the steel sheet. When the Ti content is less than 0.050%, it is difficult to obtain a desired hot-rolled steel sheet strength (tensile strength: 590 MPa or more). On the other hand, when the Ti content is excessive, the carbide containing Ti tends to be coarsened, and it becomes difficult to obtain a desired hot-rolled steel sheet strength (tensile strength: 590 MPa or more). Ti is an element that inhibits recrystallization of austenite. Therefore, when Ti is contained in a large amount exceeding 0.120%, recrystallization after finish rolling is delayed, the ferrite structure after transformation is elongated in the rolling direction, and the anisotropy of the mechanical properties of the steel sheet Becomes larger. Therefore, the Ti content is set to 0.050% or more and 0.120% or less. Preferably they are 0.060% or more and 0.110% or less, More preferably, they are 0.060% or more and 0.100% or less.
Nb:0.005%以下
本発明において、Nbは不純物(不可避的不純物)である。Nbは、Tiと同様、炭化物を形成して鋼板強度に寄与する元素である。しかしながら、Nbは、Tiよりもオーステナイトの再結晶を強く阻害する。そのため、Nb含有量が0.005%を超える場合には、仕上げ圧延後のオーステナイトの再結晶が遅延し、変態後のフェライト組織が圧延方向に伸長した組織となり、鋼板の機械的特性の異方性が大きくなる。したがって、本発明ではNb炭化物による鋼板の高強度化は行わず、Nb含有量を極力低減することとし、0.005%以下に制限する。好ましくは0.003%以下である。Nb: 0.005% or less In the present invention, Nb is an impurity (unavoidable impurity). Nb, like Ti, is an element that forms carbides and contributes to steel plate strength. However, Nb inhibits austenite recrystallization more strongly than Ti. Therefore, when the Nb content exceeds 0.005%, the recrystallization of austenite after finish rolling is delayed, and the ferrite structure after transformation becomes a structure elongated in the rolling direction. Increases sex. Therefore, in the present invention, the strength of the steel sheet is not increased by Nb carbide, and the Nb content is reduced as much as possible, and is limited to 0.005% or less. Preferably it is 0.003% or less.
B:0.0005%以下
本発明において、Bは不純物(不可避的不純物)である。Bは、Nbと同様にオーステナイトの再結晶を強く阻害し、仕上げ圧延後のオーステナイトの再結晶を遅延させる。そのため、変態後のフェライト組織が圧延方向に伸長し、鋼板の機械的特性の異方性が増大する。したがって、本発明では、B含有量を極力低減することとし、0.0005%以下に制限する。好ましくは0.0003%以下である。B: 0.0005% or less In the present invention, B is an impurity (inevitable impurity). B, like Nb, strongly inhibits the recrystallization of austenite and delays the recrystallization of austenite after finish rolling. Therefore, the ferrite structure after transformation extends in the rolling direction, and the anisotropy of the mechanical properties of the steel sheet increases. Therefore, in the present invention, the B content is reduced as much as possible, and is limited to 0.0005% or less. Preferably it is 0.0003% or less.
また、本発明において、Tiを含む炭化物の粗大化を抑制するうえでは、C、S、NおよびTiの含有量を、以下(1)式を満足するように調整することが好ましい。 Moreover, in this invention, in order to suppress the coarsening of the carbide | carbonized_material containing Ti, it is preferable to adjust content of C, S, N, and Ti so that following (1) Formula may be satisfied.
((Ti−(48/14)×N−(48/32)×S)/48)/(C/12)<1.0 ・・・ (1)
((1)式中のC、S、N、Ti:各元素の含有量(質量%))
本発明においては、フェライト相中にTiを含む炭化物を微細析出させることで所望の鋼板強度を得る。Tiを含む炭化物は、平均粒子径が極めて小さい微細炭化物となる傾向が強い。しかしながら、鋼中に含まれているTiの原子濃度がCの原子濃度以上となると、巻取り温度における固溶Ti量が急激に増大するため、Tiを含む炭化物が粗大化し易くなる。そして、Tiを含む炭化物が粗大化した結果、所望の鋼板強度(引張強さ:590MPa以上)を得ることが困難となる。以上の理由により、鋼中に含まれるCの原子%を、炭化物生成に寄与するTiの原子%よりも多くすることが望ましい。((Ti− (48/14) × N− (48/32) × S) / 48) / (C / 12) <1.0 (1)
(C, S, N, Ti in formula (1): content of each element (mass%))
In the present invention, desired steel plate strength is obtained by finely precipitating a carbide containing Ti in the ferrite phase. A carbide containing Ti has a strong tendency to become a fine carbide having an extremely small average particle diameter. However, when the atomic concentration of Ti contained in the steel is equal to or higher than the atomic concentration of C, the amount of dissolved Ti at the coiling temperature is rapidly increased, so that the carbide containing Ti is easily coarsened. As a result of coarsening of the carbide containing Ti, it becomes difficult to obtain a desired steel plate strength (tensile strength: 590 MPa or more). For the above reasons, it is desirable that the atomic percent of C contained in the steel be larger than the atomic percent of Ti that contributes to carbide formation.
後述するように、本発明においては、鋼素材に所定量のTiを添加し、熱延前の加熱で鋼素材中の炭化物を溶解し、主に熱間圧延後の巻き取り時にTiを含む炭化物を析出させる。しかしながら、鋼素材に添加したTiの全量が炭化物生成に寄与するわけではなく、鋼素材に添加したTiの一部は窒化物や硫化物の形成に使われる。巻取り温度よりも高温域では、Tiは炭化物よりも窒化物や硫化物を形成し易く、熱延鋼板の製造時、巻き取り工程の前にTiが窒化物や硫化物を形成するためである。ゆえに、鋼素材が含有するTiのうち、炭化物生成に寄与できるTi量は、「Ti−(48/14)N−(48/32)S」で表すことができる。 As will be described later, in the present invention, a predetermined amount of Ti is added to the steel material, the carbide in the steel material is dissolved by heating before hot rolling, and the carbide contains Ti mainly during winding after hot rolling. To precipitate. However, the total amount of Ti added to the steel material does not contribute to carbide formation, and a part of Ti added to the steel material is used for the formation of nitrides and sulfides. This is because Ti forms nitrides and sulfides more easily than carbides in a temperature range higher than the coiling temperature, and Ti forms nitrides and sulfides before the winding process when manufacturing a hot-rolled steel sheet. . Therefore, among the Ti contained in the steel material, the amount of Ti that can contribute to carbide generation can be represented by “Ti− (48/14) N− (48/32) S”.
以上の理由により、本発明では、Cの原子%(C/12)を、炭化物生成に寄与できるTiの原子%((Ti−(48/14)×N−(48/32)×S)/48)よりも多くする目的で、上記(1)式を満足するようにC、N、S、Tiの各元素を含有することが好ましい。より好ましくは、((Ti−(48/14)×N−(48/32)×S)/48)/(C/12)<0.9である。 For the above reasons, in the present invention, the atomic percent of C (C / 12) is converted to atomic percent of Ti that can contribute to carbide formation ((Ti− (48/14) × N− (48/32) × S) / For the purpose of increasing more than 48), it is preferable to contain each element of C, N, S, and Ti so as to satisfy the above formula (1). More preferably, ((Ti− (48/14) × N− (48/32) × S) / 48) / (C / 12) <0.9.
また、本発明の熱延鋼板は、以上の元素に加えて更に、REM、Zr、V、As、Cu、Ni、Sn、Pb、Ta、W、Mo、Cr、Sb、Mg、Ca、Co、Se、Zn、Csのうちから選ばれた1種以上を合計で1.0%以下含有してもよい。これらの元素の合計含有量が1.0%以下であれば、熱延鋼板の諸特性に悪影響を及ぼすことはない。なお、上記以外の成分は、Feおよび不可避的不純物(NbおよびBを除く不可避的不純物)である。 In addition to the above elements, the hot-rolled steel sheet of the present invention further includes REM, Zr, V, As, Cu, Ni, Sn, Pb, Ta, W, Mo, Cr, Sb, Mg, Ca, Co, One or more selected from Se, Zn, and Cs may be contained in a total of 1.0% or less. If the total content of these elements is 1.0% or less, the various properties of the hot-rolled steel sheet will not be adversely affected. Components other than the above are Fe and inevitable impurities (inevitable impurities excluding Nb and B).
次に、本発明熱延鋼板の組織の限定理由について説明する。 Next, the reason for limiting the structure of the hot-rolled steel sheet of the present invention will be described.
本発明の熱延鋼板は、上記した組成を有し、更に、フェライト相の面積率が95%以上であり、前記フェライト相の結晶粒の平均アスペクト比が3.0以下であり、前記フェライト相の結晶粒内にTiを含む炭化物が微細析出し、該炭化物の平均粒子径が10nm未満である金属組織を有する。 The hot-rolled steel sheet of the present invention has the above-described composition, and the area ratio of the ferrite phase is 95% or more, the average aspect ratio of the crystal grains of the ferrite phase is 3.0 or less, and the ferrite phase The carbide containing Ti is finely precipitated in the crystal grains, and the carbide has a metal structure having an average particle diameter of less than 10 nm.
フェライト相の面積率:95%以上
本発明においては、所望の熱延鋼板の伸びフランジ性を得るために、熱延鋼板にフェライト相を形成させることが必須となる。伸びフランジ性の向上には、熱延鋼板の金属組織がフェライト単相であることが好ましい。しかし、完全なフェライト単相で無い場合であっても、実質的にフェライト単相、すなわち、金属組織全体に対する面積率で95%以上がフェライト相であれば、上記の効果を十分に発揮する。したがって、本発明においては、熱延鋼板の金属組織を、面積率で95%以上のフェライト相を含有する組織とする。好ましくは97%以上である。In the present invention, it is essential to form a ferrite phase in the hot-rolled steel sheet in order to obtain a desired stretch-flange property of the hot-rolled steel sheet. In order to improve stretch flangeability, the metal structure of the hot-rolled steel sheet is preferably a ferrite single phase. However, even if it is not a complete ferrite single phase, the above effect is sufficiently exhibited if the ferrite single phase is substantially the ferrite phase, that is, if the area ratio with respect to the entire metal structure is 95% or more. Therefore, in the present invention, the metal structure of the hot-rolled steel sheet is a structure containing a ferrite phase having an area ratio of 95% or more. Preferably it is 97% or more.
なお、本発明の熱延鋼板において、金属組織に含有され得るフェライト相以外の組織としては、セメンタイト、パーライト、ベイナイト相、マルテンサイト相、残留オーステナイト相等が挙げられる。これらの組織が金属組織中に存在すると、鋼板の伸びフランジ性が低下する。そのため、これらの組織の面積率を低減することが好ましい。ただし、金属組織全体に対する合計面積率が5%以下であれば許容される。好ましくは3%以下である。 In the hot-rolled steel sheet of the present invention, examples of the structure other than the ferrite phase that can be contained in the metal structure include cementite, pearlite, bainite phase, martensite phase, and retained austenite phase. When these structures are present in the metal structure, the stretch flangeability of the steel sheet is lowered. Therefore, it is preferable to reduce the area ratio of these tissues. However, it is permissible if the total area ratio with respect to the entire metal structure is 5% or less. Preferably it is 3% or less.
フェライト相の結晶粒の平均アスペクト比:3.0以下
本発明におけるフェライト粒のアスペクト比は、熱延鋼板の圧延方向に平行な断面におけるフェライト粒の板厚方向の長さに対する圧延方向の長さの比である。平均アスペクト比の求め方は後述する。本発明では、アスペクト比を小さくしてフェライト粒形状を極力等軸に近づける。アスペクト比が小さければ、フェライト粒の圧延方向の長さおよび板幅方向の長さの差もまた小さくなる。そして平均アスペクト比を3.0以下にすることにより、熱延鋼板の引張強さの異方性、すなわち圧延方向および板幅方向の引張り強さの差を小さくできる。フェライト粒の平均アスペクト比が3.0を超えて大きくなると、フェライト粒形状に起因する引張強さの異方性が顕著になる。したがって、フェライト粒の平均アスペクト比は3.0以下に限定する。好ましくは2.0以下、より好ましくは1.5以下である。なお、本発明で得られた熱延鋼板のフェライト粒の平均アスペクト比は、最小で1.1であった。Average aspect ratio of ferrite phase crystal grains: 3.0 or less The aspect ratio of ferrite grains in the present invention is the length in the rolling direction relative to the length in the thickness direction of the ferrite grains in a cross section parallel to the rolling direction of the hot rolled steel sheet. Ratio. A method for obtaining the average aspect ratio will be described later. In the present invention, the aspect ratio is reduced to make the ferrite grain shape as close to the equiaxed axis as possible. If the aspect ratio is small, the difference between the length of the ferrite grains in the rolling direction and the length in the sheet width direction is also small. By setting the average aspect ratio to 3.0 or less, the anisotropy of the tensile strength of the hot-rolled steel sheet, that is, the difference in tensile strength between the rolling direction and the sheet width direction can be reduced. When the average aspect ratio of the ferrite grains increases beyond 3.0, the tensile strength anisotropy due to the ferrite grain shape becomes significant. Therefore, the average aspect ratio of ferrite grains is limited to 3.0 or less. Preferably it is 2.0 or less, More preferably, it is 1.5 or less. The average aspect ratio of the ferrite grains of the hot-rolled steel sheet obtained in the present invention was 1.1 at the minimum.
Tiを含む炭化物
本発明において熱延鋼板に微細析出させる炭化物は、Tiを含む炭化物である。熱延鋼板が炭化物構成元素としてTiのみを含有する場合、Tiを含む炭化物はTi炭化物である。また、熱延鋼板がTi以外の炭化物構成元素(V、Mo等)も含有する場合には、Ti炭化物のほかにTiとVの複合炭化物、或いは更に、Mo等の炭化物構成元素を炭化物中に含むものが挙げられる。Carbide containing Ti In the present invention, the carbide finely precipitated on the hot-rolled steel sheet is a carbide containing Ti. When the hot-rolled steel sheet contains only Ti as a carbide constituent element, the carbide containing Ti is Ti carbide. Further, when the hot-rolled steel sheet also contains carbide constituent elements (V, Mo, etc.) other than Ti, in addition to Ti carbide, composite carbide of Ti and V, or further, carbide constituent elements such as Mo, etc. in the carbide. Including.
Tiを含む炭化物の平均粒子径:10nm未満
熱延鋼板を所望の強度(引張強さ:590MPa以上)とするうえで、Tiを含む炭化物の平均粒子径は極めて重要である。そのため、本発明においてはTiを含む炭化物の平均粒子径を10nm未満とする。上記フェライト相の結晶粒内にTiを含む炭化物が微細析出すると、該炭化物が、鋼板に変形が加わった際に生じる転位の移動に対する抵抗として作用することにより熱延鋼板が高強度化される。しかしながら、Tiを含む炭化物が粗大化すると、該炭化物がまばらに析出することになり、転位を止める間隔が広がるために析出強化能は低下する。そして、Tiを含む炭化物の平均粒子径が10nm以上になると、固溶強化元素であるMn、Si含有量を減らしたことに起因する鋼板強度の低下量を補うに十分な鋼板強化能が得られない。したがって、Tiを含む炭化物の平均粒子径は10nm未満とする。より好ましくは6nm以下である。Average particle diameter of carbides containing Ti: less than 10 nm The average particle diameter of carbides containing Ti is extremely important for making a hot-rolled steel sheet have a desired strength (tensile strength: 590 MPa or more). Therefore, in this invention, the average particle diameter of the carbide | carbonized_material containing Ti shall be less than 10 nm. When carbide containing Ti is finely precipitated in the ferrite phase crystal grains, the carbide acts as a resistance against dislocation movement that occurs when the steel sheet is deformed, thereby increasing the strength of the hot-rolled steel sheet. However, when the carbide containing Ti is coarsened, the carbide is sparsely precipitated, and the interval for stopping the dislocation is widened, so that the precipitation strengthening ability is lowered. And when the average particle diameter of the carbide containing Ti is 10 nm or more, sufficient steel plate strengthening ability to compensate for the decrease in steel plate strength caused by reducing the content of Mn and Si as solid solution strengthening elements is obtained. Absent. Therefore, the average particle diameter of the carbide containing Ti is less than 10 nm. More preferably, it is 6 nm or less.
なお、特に発明の効果を限定するものではないが、本発明における微細析出物(Tiを含む炭化物)が、観察する角度によっては列状に並んでいるように観察される場合がある。しかし、この場合でも析出物の列が観察された平面内では、実際には析出物がランダムに分布しており、透過型電子顕微鏡で観察すると、析出物が列状に観察されない場合が多い。 In addition, although the effect of invention is not specifically limited, the fine precipitate (carbide containing Ti) in this invention may be observed so that it may be located in a line depending on the angle to observe. However, even in this case, the precipitates are actually randomly distributed in the plane where the rows of precipitates are observed, and when observed with a transmission electron microscope, the precipitates are often not observed in rows.
以上のように組成と組織を規定することで、所望の強度(引張強さ:590MPa以上)を有し、伸びフランジ性に優れ、且つ、引張強さの異方性の小さい高強度熱延鋼板が得られる。また、鋼板に耐食性を付与する目的で、本発明熱延鋼板の表面にめっき層を設けても、上記した本発明の効果を損なうことはない。なお、本発明において鋼板表面に設けるめっき層の種類は特に限定されず、電気めっき層、溶融めっき層のいずれも適用可能である。めっき層の合金成分も特に問わず、亜鉛めっき層、合金化亜鉛めっき層などが好適な例として挙げられる。勿論、本発明はこれらの例に限定されない。表面にめっき層を形成することにより、熱延鋼板の耐食性が向上し、厳しい腐食環境下で使用される自動車部品などへの適用が可能になる。 By defining the composition and structure as described above, a high-strength hot-rolled steel sheet having a desired strength (tensile strength: 590 MPa or more), excellent stretch flangeability, and small tensile strength anisotropy. Is obtained. Moreover, even if a plated layer is provided on the surface of the hot-rolled steel sheet of the present invention for the purpose of imparting corrosion resistance to the steel sheet, the above-described effects of the present invention are not impaired. In the present invention, the type of the plating layer provided on the surface of the steel sheet is not particularly limited, and any of an electroplating layer and a hot dipping layer can be applied. The alloy component of the plating layer is not particularly limited, and preferred examples include a galvanized layer and an alloyed galvanized layer. Of course, the present invention is not limited to these examples. By forming a plating layer on the surface, the corrosion resistance of the hot-rolled steel sheet is improved, and it becomes possible to apply it to automobile parts used in severe corrosive environments.
次に、本発明の熱延鋼板の製造方法について説明する。 Next, the manufacturing method of the hot rolled steel sheet of the present invention will be described.
本発明は、上記した組成の鋼素材を、オーステナイト単相域に加熱し、熱間圧延を施した後、冷却し、巻き取り、熱延鋼板とする。この際、前記熱間圧延の仕上げ圧延において、900℃以上1100℃以下の温度域で圧下率15%以上の圧延を3回以上行い、仕上げ圧延温度を860℃以上とし、前記仕上げ圧延終了後、850℃以上の温度域で0.3s以上保持した後前記冷却を開始し、前記冷却の850℃から750℃までの平均冷却速度を30℃/s以上とし、前記巻き取りの巻取り温度を580℃以上750℃以下とすることを特徴とする。 In the present invention, a steel material having the above composition is heated to an austenite single-phase region, subjected to hot rolling, then cooled, wound, and formed into a hot-rolled steel sheet. At this time, in the finish rolling of the hot rolling, rolling at a reduction rate of 15% or more is performed three times or more in a temperature range of 900 ° C. or more and 1100 ° C. or less, the finish rolling temperature is 860 ° C. or more, and after the finish rolling is finished, The cooling is started after being held for 0.3 s or more in a temperature range of 850 ° C. or more, the average cooling rate from 850 ° C. to 750 ° C. of the cooling is 30 ° C./s or more, and the winding temperature of the winding is 580 It is characterized by being not less than 750 ° C. and not more than 750 ° C.
本発明において、鋼の溶製方法は特に限定されず、転炉、電気炉等、公知の溶製方法を採用することができる。また、真空脱ガス炉にて2次精錬を行ってもよい。その後、溶鋼からスラブを鋳造する。生産性等の問題から連続鋳造法によりスラブ(鋼素材)とするのが好ましい。但し、造塊−分塊圧延法、薄スラブ連鋳法等、公知の鋳造方法でスラブとしてもよい。なお、本発明では、鋼板の加工性(伸びフランジ性等)の向上を目的として偏析の原因となるMnやSiの含有量を抑制している。そのため、偏析の抑制に有利な連続鋳造法を採用すると、本発明の効果がより一層顕著となる。 In the present invention, the method for melting steel is not particularly limited, and a known melting method such as a converter or an electric furnace can be employed. Further, secondary refining may be performed in a vacuum degassing furnace. Thereafter, a slab is cast from the molten steel. From the viewpoint of productivity and the like, the slab (steel material) is preferably formed by a continuous casting method. However, the slab may be formed by a known casting method such as ingot-bundling rolling or continuous slab casting. In the present invention, the content of Mn and Si causing segregation is suppressed for the purpose of improving the workability (stretch flangeability, etc.) of the steel sheet. Therefore, when the continuous casting method advantageous for suppressing segregation is employed, the effect of the present invention becomes more remarkable.
上記の如く得られた鋼素材に、熱間圧延を施すが、本発明においては、熱間圧延に先立ち、鋼素材をオーステナイト単相域に加熱する。熱間圧延前の鋼素材がオーステナイト単相域に加熱されていないと、鋼素材中に存在する炭化物(Tiを含む炭化物)の再溶解が進行せず、圧延後にTiを含む炭化物の微細析出が実現できない。 The steel material obtained as described above is hot-rolled. In the present invention, the steel material is heated to an austenite single-phase region prior to hot rolling. If the steel material before hot rolling is not heated to the austenite single-phase region, remelting of carbides (carbides containing Ti) in the steel materials will not proceed, and fine precipitation of carbides containing Ti after rolling will not occur. Cannot be realized.
以上の理由により、本発明では熱間圧延に先立ち、鋼素材をオーステナイト単相域、好ましくは1200℃以上に加熱する。但し、鋼素材の加熱温度が高くなりすぎると、鋼素材の表面が過度に酸化され、表面近傍にTiO2が生じる。そして、TiO2が生じる結果、表面近傍において、炭化物を形成するためのTiが少なくなり、その結果最終的に得られる鋼板の表面近傍の硬さが低下し易くなる。したがって、上記加熱温度は1350℃以下とすることが好ましい。なお、鋼素材に熱間圧延を施すに際し、鋳造後の鋼素材(スラブ)がオーステナイト単相域の温度となっている場合には、鋼素材を加熱することなく、或いは短時間加熱後、直送圧延してもよい。For the above reasons, in the present invention, prior to hot rolling, the steel material is heated to an austenite single phase region, preferably 1200 ° C. or higher. However, if the heating temperature of the steel material becomes too high, the surface of the steel material is excessively oxidized, and TiO 2 is generated in the vicinity of the surface. As a result of the generation of TiO 2 , Ti for forming carbide is reduced in the vicinity of the surface, and as a result, the hardness in the vicinity of the surface of the finally obtained steel sheet is likely to decrease. Therefore, the heating temperature is preferably 1350 ° C. or lower. In addition, when hot rolling the steel material, if the steel material (slab) after casting has a temperature in the austenite single-phase region, the steel material is not heated or directly heated after being heated for a short time. You may roll.
熱間圧延は、通常、粗圧延と仕上げ圧延からなる。本発明において粗圧延の条件は特に限定されない。また、特に薄スラブ鋳造法を採用した場合には、粗圧延を省略してもよい。仕上げ圧延は、以下の条件で行う。 Hot rolling usually consists of rough rolling and finish rolling. In the present invention, the conditions for rough rolling are not particularly limited. In particular, when a thin slab casting method is employed, rough rolling may be omitted. Finish rolling is performed under the following conditions.
900℃以上1100℃以下の温度域での圧延:圧下率15%以上の圧延を3回以上
一般に、熱間圧延によって引き伸ばされたオーステナイトから変態したフェライトは、圧延方向に伸長した組織を形成し易いため、鋼板の引張強さの異方性が大きくなる要因となる。一方、熱間圧延後再結晶したオーステナイトは、等軸となるため、再結晶後のオーステナイトから変態したフェライトが等軸組織となり易く、鋼板の引張強さの異方性が小さくなる。このため、本発明では、仕上げ圧延後の再結晶を促進させ、再結晶後のオーステナイトからフェライト変態させる必要がある。Rolling in a temperature range of 900 ° C. or higher and 1100 ° C. or lower: rolling at a rolling reduction of 15% or more 3 times or more Generally, ferrite transformed from austenite stretched by hot rolling tends to form a structure elongated in the rolling direction. Therefore, the anisotropy of the tensile strength of the steel sheet becomes a factor. On the other hand, since austenite recrystallized after hot rolling becomes equiaxed, the ferrite transformed from the austenite after recrystallization tends to have an equiaxed structure, and the anisotropy of the tensile strength of the steel sheet becomes small. For this reason, in the present invention, it is necessary to promote recrystallization after finish rolling and to transform ferrite from austenite after recrystallization.
ここで、適切な温度域において圧延によって加工ひずみを導入することにより、オーステナイトを再結晶させることができる。しかしながら、オーステナイト組織が粗大なうちは、鋼中に一部再結晶しない領域が残存する。この未再結晶部は、オーステナイト組織が圧延方向に大きく伸長しており、また、後述する仕上げ圧延後の保持において再結晶を起こし難い領域であるため、異方性を誘発する。この未再結晶オーステナイト部は、圧延・再結晶を繰り返すことで徐々に消滅し、これによって均一で細粒な再結晶オーステナイト組織とすることができる。オーステナイトの再結晶は900℃以上1100℃以下の温度域において圧下率15%以上の圧延を施すことで達成される。そして、鋼板全体に亘り均一で細粒な再結晶オーステナイト組織を得るためには、900℃以上1100℃以下の温度域において圧下率15%以上の圧延を最低3回は行う必要がある。好ましくは4回以上である。なお、仕上げ圧延スタンド数の制約から、上限は実質7回程度である。 Here, austenite can be recrystallized by introducing processing strain by rolling in an appropriate temperature range. However, as long as the austenite structure is coarse, a region that does not partially recrystallize remains in the steel. This non-recrystallized portion has an austenite structure that extends greatly in the rolling direction, and also induces anisotropy because it is a region in which recrystallization hardly occurs during holding after finish rolling described later. This non-recrystallized austenite part disappears gradually by repeating rolling and recrystallization, and thereby a uniform and fine-grained recrystallized austenite structure can be obtained. Austenite recrystallization is achieved by rolling at a rolling reduction of 15% or more in a temperature range of 900 ° C. or higher and 1100 ° C. or lower. In order to obtain a recrystallized austenite structure that is uniform and fine throughout the entire steel plate, it is necessary to perform rolling at a reduction rate of 15% or more at least three times in a temperature range of 900 ° C. or higher and 1100 ° C. or lower. Preferably it is 4 times or more. Note that the upper limit is substantially about seven due to the restriction on the number of finish rolling stands.
仕上げ圧延温度:860℃以上
前述のように、引張強さの異方性の小さい鋼板とするためには、仕上げ圧延後の再結晶を促進させ、再結晶後のオーステナイトからフェライト変態させる必要がある。ここで、仕上げ圧延温度が860℃未満であると、仕上げ圧延後、加速冷却を開始までの間に、再結晶によるオーステナイト粒の等軸化が進行し難くなる。それゆえ、仕上げ圧延温度が860℃未満であると、変態後のフェライト粒の形状が圧延方向に伸長した形状となり、鋼板の引張強さの異方性が大きくなり易い。以上の理由により、仕上げ圧延温度は860℃以上とする。また、880℃以上とすることが好ましく、900℃以上とすることがより好ましい。但し、仕上げ圧延温度が必要以上に高くなると、表面の2次スケールによる疵や荒れが生じ易くなるため、仕上げ圧延温度の実質的な上限は1050℃程度である。Finish rolling temperature: 860 ° C. or higher As described above, in order to obtain a steel sheet with low tensile strength anisotropy, it is necessary to promote recrystallization after finish rolling and to transform ferrite from austenite after recrystallization. . Here, when the finish rolling temperature is less than 860 ° C., it becomes difficult for the austenite grains to be equiaxed by recrystallization before the start of accelerated cooling after finish rolling. Therefore, if the finish rolling temperature is less than 860 ° C., the shape of the ferrite grains after transformation becomes a shape elongated in the rolling direction, and the anisotropy of the tensile strength of the steel sheet tends to increase. For the above reasons, the finish rolling temperature is set to 860 ° C. or higher. Moreover, it is preferable to set it as 880 degreeC or more, and it is more preferable to set it as 900 degreeC or more. However, if the finish rolling temperature is higher than necessary, wrinkles and roughness due to the secondary scale of the surface are likely to occur, so the substantial upper limit of the finish rolling temperature is about 1050 ° C.
仕上げ圧延終了後、850℃以上の温度域での保持時間:0.3s以上
本発明では、仕上げ圧延直後に加速冷却を開始せず、高温域に所定時間保持することで、仕上げ圧延によって伸長したオーステナイトの再結晶を進行させ、変態後のフェライト粒を等軸または等軸に近い形状にする。オーステナイトの再結晶を十分に進行させるには、仕上げ圧延後の鋼板を850℃以上の温度域で0.3秒以上保持する必要がある。好ましくは0.5秒以上、より好ましくは0.8秒以上である。After completion of finish rolling, holding time in a temperature range of 850 ° C. or higher: 0.3 s or more In the present invention, accelerated cooling is not started immediately after finish rolling, but is held in a high temperature range for a predetermined time, thereby extending by finish rolling. Austenite recrystallization proceeds to make the transformed ferrite grains equiaxed or close to equiaxed. In order to sufficiently advance the recrystallization of austenite, it is necessary to hold the steel sheet after finish rolling for 0.3 seconds or more in a temperature range of 850 ° C. or more. Preferably it is 0.5 second or more, More preferably, it is 0.8 second or more.
但し、上記温度域での保持時間があまりに長くなると、冷却設備の長さの制約から、目標の巻取り温度まで冷却することが困難となる。したがって、上記温度域での保持時間の実質的な上限は10秒程度である。仕上げ圧延終了後、850℃以上の温度域で0.3秒以上保持した後、加速冷却を開始する。加速冷却は、以下の条件で行う。 However, if the holding time in the above temperature range becomes too long, it becomes difficult to cool to the target winding temperature due to the restriction of the length of the cooling equipment. Therefore, the substantial upper limit of the holding time in the temperature range is about 10 seconds. After completion of finish rolling, accelerated cooling is started after holding at a temperature range of 850 ° C. or higher for 0.3 seconds or longer. Accelerated cooling is performed under the following conditions.
850℃から750℃までの平均冷却速度:30℃/s以上
平均粒子径が10nm未満の微細な炭化物(Tiを含む炭化物)を析出させるためには、加速冷却し、可能な限り低い温度でγ→α変態が生じるようにすることが必要である。850℃から750℃までの温度域における平均冷却速度が30℃/s未満になると、γ→α変態が高温で生じるようになり、フェライト中に析出した炭化物が粗大化し易くなる。したがって、850℃から750℃までの温度域における平均冷却速度は30℃/s以上とする。好ましくは50℃/s以上である。但し、上記温度域における平均冷却速度が過剰に大きくなると、巻取り温度の制御が困難となり安定した強度が得られ難くなるおそれがあるため、上記温度域における平均冷却速度は300℃/s以下とすることが好ましい。Average cooling rate from 850 ° C. to 750 ° C .: 30 ° C./s or more In order to precipitate fine carbides (carbides containing Ti) having an average particle diameter of less than 10 nm, accelerated cooling and γ at the lowest possible temperature → It is necessary to make α transformation occur. When the average cooling rate in the temperature range from 850 ° C. to 750 ° C. is less than 30 ° C./s, the γ → α transformation occurs at a high temperature, and the carbides precipitated in the ferrite are easily coarsened. Therefore, the average cooling rate in the temperature range from 850 ° C. to 750 ° C. is set to 30 ° C./s or more. Preferably it is 50 degrees C / s or more. However, if the average cooling rate in the temperature range becomes excessively large, it is difficult to control the coiling temperature and it may be difficult to obtain a stable strength. Therefore, the average cooling rate in the temperature range is 300 ° C./s or less. It is preferable to do.
巻取り温度:580℃以上750℃以下
巻き取り温度の適正化は、フェライト中に微細な炭化物(Tiを含む炭化物)を析出させ、且つ、熱延鋼板を所望の金属組織(フェライト相の面積率:95%以上、Tiを含む炭化物の平均粒子径:10nm未満)とするうえで重要である。巻取り温度が580℃未満であると、マルテンサイトやベイナイトが生じ易くなり、金属組織を実質的にフェライト単相組織とすることが困難となる。一方、巻取り温度が750℃を超えると、巻き取り後のコイルの冷却中にTiを含む炭化物の粗大化が促進されるため、Tiを含む炭化物の平均粒子径を10nm未満とすることが困難となる。したがって、巻取り温度は580℃以上750℃以下とする。好ましくは600℃以上680℃以下である。Winding temperature: 580 ° C. or higher and 750 ° C. or lower The optimization of the winding temperature is performed by precipitating fine carbides (carbides containing Ti) in the ferrite and forming a desired metal structure (area ratio of ferrite phase) in the hot rolled steel sheet. : 95% or more, average particle diameter of carbide containing Ti: less than 10 nm). When the coiling temperature is less than 580 ° C., martensite and bainite are likely to be generated, and it becomes difficult to make the metal structure substantially a ferrite single phase structure. On the other hand, when the coiling temperature exceeds 750 ° C., the coarsening of the carbide containing Ti is promoted during cooling of the coil after winding, and therefore it is difficult to make the average particle diameter of the carbide containing Ti less than 10 nm. It becomes. Therefore, the coiling temperature is set to 580 ° C. or higher and 750 ° C. or lower. Preferably they are 600 degreeC or more and 680 degrees C or less.
更に、巻き取り後のコイルを580℃以上750℃以下の温度範囲に60s以上保持すると、均一な組織が得られ易くなるため、一層好ましい。なお、本発明の熱延鋼板の板厚は、特に限定されないが、概ね1.0mm以上8.0mm以下である。 Furthermore, holding the coil after winding in a temperature range of 580 ° C. or higher and 750 ° C. or lower for 60 seconds or more is more preferable because a uniform structure can be easily obtained. The thickness of the hot-rolled steel sheet of the present invention is not particularly limited, but is generally 1.0 mm or more and 8.0 mm or less.
以上により、引張強さ(TS):590MPa以上の高強度と優れた加工性(特に伸びフランジ性)を有し、且つ引張強さの異方性が小さい高強度熱延鋼板が得られる。なお、本発明においては、黒皮材(熱延まま材)、白皮材(熱延酸洗材)のいずれの場合であっても、上記所望の効果を発現する。 As described above, a high-strength hot-rolled steel sheet having high tensile strength (TS): 590 MPa or more, excellent workability (particularly stretch flangeability) and small tensile strength anisotropy can be obtained. In the present invention, the desired effect is exhibited regardless of whether the material is a black skin material (as hot rolled material) or a white skin material (hot rolled pickling material).
また、本発明においては、以上のようにして製造された熱延鋼板に対し、めっき処理を施して鋼板表面にめっき層を形成してもよい。めっき層を形成しても、本発明の効果を損なうことはない。めっき処理は、電気めっき、溶融めっきのいずれも適用可能である。また、めっき層の合金成分も特に問わず、溶融亜鉛めっき層、合金化溶融亜鉛めっき層などが好適な例として挙げられる。勿論、本発明はこれらの例に限定されない。アルミもしくはアルミ合金等をめっきすることもできる。 In the present invention, the hot-rolled steel sheet produced as described above may be plated to form a plating layer on the steel sheet surface. Even if the plating layer is formed, the effect of the present invention is not impaired. As the plating treatment, either electroplating or hot dipping can be applied. Further, the alloy component of the plating layer is not particularly limited, and preferred examples include a hot dip galvanized layer and an alloyed hot dip galvanized layer. Of course, the present invention is not limited to these examples. Aluminum or aluminum alloy can also be plated.
本発明により得られる熱延鋼板は、常温で行われるプレス成形用素材に好適なほか、プレス前の鋼板を400℃から750℃に加温した後直ちにプレス成形する温間成形にも好適である。 The hot-rolled steel sheet obtained by the present invention is suitable for a material for press molding performed at room temperature, and also suitable for warm forming in which a steel sheet before pressing is heated from 400 ° C. to 750 ° C. and then press-formed immediately. .
溶鋼を通常公知の手法により溶製、連続鋳造して、表1に示す組成を有する肉厚300mmのスラブ(鋼素材)とした。これらのスラブを、表2に示す条件で加熱し、粗圧延後、表2に示す条件で仕上げ圧延を施し、仕上げ圧延終了後、850℃以上の温度域に所定時間保持し、加速冷却し、巻き取り、板厚:2.3mmの熱延鋼板とした。 Molten steel was melted and continuously cast by a generally known method to obtain a slab (steel material) having a thickness of 300 mm having the composition shown in Table 1. These slabs are heated under the conditions shown in Table 2, and after rough rolling, finish rolling is performed under the conditions shown in Table 2. After finishing rolling, the slab is kept in a temperature range of 850 ° C. or higher for a predetermined time, accelerated cooling, Winding, plate thickness: 2.3 mm hot rolled steel sheet.
続いて、上記により得られた熱延鋼板を酸洗して表層スケールを除去した後、一部の熱延鋼板(鋼板No.S40、S41、S42)については焼鈍温度720℃の溶融亜鉛めっきラインに通板し、480℃の亜鉛めっき浴(めっき組成:0.1mass%Al−Zn)中に浸漬し、片面当たり付着量45g/m2の溶融亜鉛めっき層を鋼板の表面に形成して溶融亜鉛めっき鋼板(GI材)とした。また、更に一部の熱延鋼板(No.S43、S44)については、上記の如く溶融亜鉛めっき層を形成したのち、520℃で合金化処理を施して合金化溶融亜鉛めっき鋼板(GA材)とした。Subsequently, after pickling the hot-rolled steel sheet obtained above to remove the surface scale, some hot-rolled steel sheets (steel plates No. S40, S41, S42) are hot-dip galvanized lines with an annealing temperature of 720 ° C. And immersed in a 480 ° C. zinc plating bath (plating composition: 0.1 mass% Al—Zn) to form a hot dip galvanized layer with an adhesion amount of 45 g / m 2 per side on the surface of the steel plate and melt A galvanized steel sheet (GI material) was used. Further, for some hot-rolled steel sheets (No. S43, S44), after forming a hot-dip galvanized layer as described above, alloying treatment is performed at 520 ° C. to form an alloyed hot-dip galvanized steel sheet (GA material). It was.
上記により得られた熱延鋼板(熱延鋼板、GI材、GA材)から試験片を採取し、組織観察、引張試験、穴拡げ試験を行い、フェライト相の面積率、フェライト相以外の組織の種類および面積率、フェライト相の結晶粒の平均アスペクト比、Tiを含む炭化物の平均粒子径、引張強さ、伸び、穴拡げ率(伸びフランジ性)を求めた。観察方法、試験方法は次のとおりとした。
(i)組織観察
得られた熱延鋼板(熱延鋼板、GI材、GA材)から試験片を採取し、試験片の圧延方向と平行な板厚1/4位置の断面(L断面)を研磨し、ナイタールで腐食した後、光学顕微鏡(倍率:400倍)および走査型電子顕微鏡(倍率:2000倍)にて組織写真を撮影した。次いで、撮影した組織写真を用い、画像解析装置によりフェライト相、フェライト相以外の組織の種類、および、それらの面積率を求めた。 Samples are taken from the hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material) obtained as described above, and subjected to structure observation, tensile test, and hole expansion test. The type and area ratio, the average aspect ratio of the crystal grains of the ferrite phase, the average particle diameter of the carbide containing Ti, the tensile strength, the elongation, and the hole expansion ratio (stretch flangeability) were determined. The observation method and test method were as follows.
(I) Microstructure observation A test piece is taken from the obtained hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material), and a cross-section (L cross-section) at a 1/4 thickness position parallel to the rolling direction of the test piece is obtained. After polishing and corroding with nital, a structure photograph was taken with an optical microscope (magnification: 400 times) and a scanning electron microscope (magnification: 2000 times). Subsequently, using the photographed structure | tissue photograph, the kind of structure | tissue other than a ferrite phase and a ferrite phase and those area ratios were calculated | required with the image-analysis apparatus.
また、得られた熱延鋼板(熱延鋼板、GI材、GA材)から試験片を採取し、試験片の圧延方向と平行な板厚1/4位置の断面(L断面)を研磨し、ナイタールで腐食した後、走査型電子顕微鏡(倍率:800倍)にて組織写真を撮影した。次いで、撮影した組織写真を用い、画像解析によりフェライト相の結晶粒の平均アスペクト比を求めた。ここで、フェライト相の結晶粒の平均アスペクト比は、800倍で撮影した写真について水平線および垂直線をそれぞれ3本引いたときの、垂直線がフェライト結晶粒を切る長さ(平均結晶粒切片長さ)に対する、水平線の平均結晶粒切片長さの比で求めた。なお、前記写真には、圧延方向と平行な方向に水平線を引くこととする。 In addition, a test piece is taken from the obtained hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material), and a cross section (L cross section) at a 1/4 thickness position parallel to the rolling direction of the test piece is polished. After corroding with nital, a structure photograph was taken with a scanning electron microscope (magnification: 800 times). Next, the average aspect ratio of the ferrite phase crystal grains was determined by image analysis using the photographed structure photograph. Here, the average aspect ratio of the ferrite phase crystal grains is the length of the vertical lines cutting the ferrite crystal grains when three horizontal lines and three vertical lines are drawn for each photograph taken at a magnification of 800 times (average crystal grain section length). )) To the average grain section length of the horizontal line. In the photograph, a horizontal line is drawn in a direction parallel to the rolling direction.
また、得られた熱延鋼板(熱延鋼板、GI材、GA材)の板厚1/4位置から作製した薄膜を透過型電子顕微鏡(TEM)によって観察し、Tiを含む炭化物の平均粒子径を求めた。Tiを含む炭化物の平均粒子径は、透過型電子顕微鏡(倍率:340000倍)にて撮影した写真を用い、5視野合計で最低100個の炭化物(Tiを含む炭化物)について粒子径を測定し、その平均値を平均粒子径とした。なお、炭化物の粒子径は、炭化物の最大径d(最も大きい部分の直径)とそれに直交する方向の径(厚さ)tとを測定し、これらの算術平均値ddef=(d+t)/2として求めた。また、Tiを含む炭化物の同定は、TEMに付帯するEDXによって分析した。
(ii)引張試験
得られた熱延鋼板(熱延鋼板、GI材、GA材)から、圧延方向に対して直角方向(C方向)、圧延方向と平行方向(L方向)、圧延方向と45度方向(D方向)を引張方向とするJIS5号引張試験片(JIS Z 2241)を採取し、JIS Z 2241の規定に準拠した引張試験を行い、引張強さ(TS)、伸び(El)を測定した。次いで、C方向、L方向、D方向のうち、最も引張強さが高い方向と最も引張強さが低い方向の引張強さの差ΔTSを求め、スプリングバック量変動をもたらす引張強さの異方性の指標とした。なお、C方向の引張強さが590MPa以上である場合を、所望の鋼板強度を有するものと評価した。また、ΔTSが50MPa未満である場合を、引張強さの異方性が小さいと評価した。
(iii)穴拡げ試験
得られた熱延鋼板(熱延鋼板、GI材、GA材)から、試験片(大きさ:130mm×130mm)を採取し、該試験片に初期直径d0:10mmφの穴を打抜き加工(クリアランス:試験片板厚の12.5%)で形成した。これら試験片を用いて、穴拡げ試験を実施した。すなわち、該穴に打ち抜き時のポンチ側から頂角:60°の円錐ポンチを挿入し、該穴を押し広げ、亀裂が鋼板(試験片)の板厚方向に貫通したときの穴の径dを測定し、次式で穴拡げ率λ(%)を算出した。ここで、穴拡げ率λ:100%以上である場合を、加工性が優れていると判断する。Moreover, the thin film produced from the plate | board thickness 1/4 position of the obtained hot-rolled steel plate (hot-rolled steel plate, GI material, GA material) is observed with a transmission electron microscope (TEM), and the average particle diameter of the carbide containing Ti Asked. The average particle diameter of the carbide containing Ti is measured using a photograph taken with a transmission electron microscope (magnification: 340000 times) for a minimum of 100 carbides (carbide containing Ti) in a total of five fields of view, The average value was defined as the average particle size. The carbide particle diameter is determined by measuring the maximum diameter d of the carbide (the diameter of the largest portion) and the diameter (thickness) t in the direction orthogonal thereto, and the arithmetic average value d def = (d + t) / 2. As sought. Moreover, the carbide | carbonized_material containing Ti was analyzed by EDX attached to TEM.
(Ii) Tensile test From the obtained hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material), the direction perpendicular to the rolling direction (C direction), the rolling direction and parallel direction (L direction), the rolling direction and 45 A JIS No. 5 tensile test piece (JIS Z 2241) with the tensile direction (D direction) as the tensile direction is sampled and subjected to a tensile test in accordance with the provisions of JIS Z 2241 to determine the tensile strength (TS) and elongation (El). It was measured. Next, the difference in tensile strength ΔTS between the direction with the highest tensile strength and the direction with the lowest tensile strength among the C direction, L direction, and D direction is obtained, and the anisotropic tensile strength that causes fluctuations in the springback amount. It was used as an index of sex. In addition, the case where the tensile strength of a C direction was 590 Mpa or more was evaluated as having the desired steel plate strength. Moreover, when ΔTS was less than 50 MPa, it was evaluated that the anisotropy of tensile strength was small.
(Iii) Hole expansion test From the obtained hot-rolled steel sheet (hot-rolled steel sheet, GI material, GA material), a test piece (size: 130 mm × 130 mm) was sampled, and an
穴拡げ率λ(%)={(d−d0)/d0}×100
以上により得られた結果を、表3に示す。また、上記(i)により得られたフェライト相の結晶粒の平均アスペクト比と、上記(ii)により得られたΔTSとの関係を、図1に示す。Hole expansion rate λ (%) = {(d−d 0 ) / d 0 } × 100
The results obtained as described above are shown in Table 3. Moreover, the relationship between the average aspect ratio of the crystal grains of the ferrite phase obtained by the above (i) and ΔTS obtained by the above (ii) is shown in FIG.
表3に示すように、発明例はいずれも、引張強さTS:590MPa以上の高強度と、伸びEl:22%以上であり且つ、穴拡げ率λ:100%以上の優れた加工性を兼備した熱延鋼板が得られている。また、発明例の熱延鋼板はいずれも、フェライト粒の平均アスペクト比が3.0以下であり、ΔTSが50MPa未満とTS異方性が抑制されている。更に、図1に示すように、フェライト粒の平均アスペクト比を3.0以下とすることにより、ΔTSが50MPa未満となり、TSの異方性の小さい熱延鋼板が得られることが理解できる。一方、本発明の範囲を外れる比較例の熱延鋼板は、所定の高強度が確保できていないか、十分な穴拡げ率が確保できていない、またはフェライト粒の平均アスペクト比が3.0を超え、ΔTSが50MPa以上となっている。
As shown in Table 3, all the inventive examples have high tensile strength TS: 590 MPa or more, elongation El: 22% or more, and excellent workability of hole expansion ratio λ: 100% or more. A hot-rolled steel sheet is obtained. Moreover, as for the hot-rolled steel plate of the example of an invention, the average aspect-ratio of a ferrite grain is 3.0 or less, (DELTA) TS is less than 50 Mpa, and TS anisotropy is suppressed. Furthermore, as shown in FIG. 1, it can be understood that by setting the average aspect ratio of ferrite grains to 3.0 or less, ΔTS is less than 50 MPa, and a hot-rolled steel sheet having a small TS anisotropy can be obtained. On the other hand, the hot-rolled steel sheet of the comparative example outside the scope of the present invention does not ensure a predetermined high strength, does not secure a sufficient hole expansion rate, or has an average aspect ratio of ferrite grains of 3.0. And ΔTS is 50 MPa or more.
Claims (5)
フェライト相の面積率が95%以上であり、前記フェライト相の結晶粒の平均アスペクト比が3.0以下であり、前記フェライト相の結晶粒内にTiを含む炭化物が微細析出し、該炭化物の平均粒子径が10nm未満である組織を有し、
引張強さが590MPa以上であることを特徴とする高強度熱延鋼板。In mass%, C: more than 0.010% and 0.070% or less, Si: 0.30% or less, Mn: more than 0.30% and 1.00% or less, P: 0.030% or less, S: 0.00. Nb and B containing 030% or less, Al: 0.10% or less, N: 0.0100% or less, Ti: 0.050% or more and 0.120% or less, and Nb and B as impurity elements in Nb : 0.005% or less, B: 0.0005% or less, the balance is composed of Fe and inevitable impurities,
The area ratio of the ferrite phase is 95% or more, the average aspect ratio of the ferrite phase crystal grains is 3.0 or less, and a carbide containing Ti is finely precipitated in the ferrite phase crystal grains. Having a structure with an average particle size of less than 10 nm,
A high-strength hot-rolled steel sheet having a tensile strength of 590 MPa or more.
記
((Ti−(48/14)×N−(48/32)×S)/48)/(C/12)<1.0 ・・・(1)
((1)式中のC、S、N、Ti:各元素の含有量(質量%))The high-strength hot-rolled steel sheet according to claim 1, wherein the composition satisfies the following formula (1).
((Ti− (48/14) × N− (48/32) × S) / 48) / (C / 12) <1.0 (1)
(C, S, N, Ti in formula (1): content of each element (mass%))
Applications Claiming Priority (3)
Application Number | Priority Date | Filing Date | Title |
---|---|---|---|
JP2014020534 | 2014-02-05 | ||
JP2014020534 | 2014-02-05 | ||
PCT/JP2015/000499 WO2015118864A1 (en) | 2014-02-05 | 2015-02-04 | High-strength hot-rolled steel sheet and production method therefor |
Publications (2)
Publication Number | Publication Date |
---|---|
JPWO2015118864A1 true JPWO2015118864A1 (en) | 2017-03-23 |
JP6224704B2 JP6224704B2 (en) | 2017-11-01 |
Family
ID=53777676
Family Applications (1)
Application Number | Title | Priority Date | Filing Date |
---|---|---|---|
JP2015521173A Active JP6224704B2 (en) | 2014-02-05 | 2015-02-04 | Manufacturing method of high strength hot-rolled steel sheet |
Country Status (2)
Country | Link |
---|---|
JP (1) | JP6224704B2 (en) |
WO (1) | WO2015118864A1 (en) |
Families Citing this family (4)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JP6135759B2 (en) * | 2014-02-05 | 2017-05-31 | Jfeスチール株式会社 | High strength hot rolled steel sheet and method for producing the same |
MX2018001946A (en) * | 2015-08-19 | 2018-06-19 | Jfe Steel Corp | High-strength steel sheet and production method for same. |
MX2018001947A (en) * | 2015-08-19 | 2018-06-19 | Jfe Steel Corp | High-strength steel sheet and production method for same. |
KR101899674B1 (en) * | 2016-12-19 | 2018-09-17 | 주식회사 포스코 | High strength steel sheet having excellent burring property in low-temperature region and manufacturing method for same |
Citations (4)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JP2009235440A (en) * | 2008-03-26 | 2009-10-15 | Jfe Steel Corp | High-yield ratio and high-strength cold rolled steel sheet |
WO2013069210A1 (en) * | 2011-11-08 | 2013-05-16 | Jfeスチール株式会社 | High-tensile-strength hot-rolled plated steel sheet and method for producing same |
JP2013133499A (en) * | 2011-12-27 | 2013-07-08 | Jfe Steel Corp | High-tensile hot-rolled steel sheet excellent in bendability and manufacturing method therefor |
JP2013133525A (en) * | 2011-12-27 | 2013-07-08 | Jfe Steel Corp | High-tensile hot-rolled steel sheet excellent in blanking property and stretch flange formability, and manufacturing method therefor |
-
2015
- 2015-02-04 JP JP2015521173A patent/JP6224704B2/en active Active
- 2015-02-04 WO PCT/JP2015/000499 patent/WO2015118864A1/en active Application Filing
Patent Citations (4)
Publication number | Priority date | Publication date | Assignee | Title |
---|---|---|---|---|
JP2009235440A (en) * | 2008-03-26 | 2009-10-15 | Jfe Steel Corp | High-yield ratio and high-strength cold rolled steel sheet |
WO2013069210A1 (en) * | 2011-11-08 | 2013-05-16 | Jfeスチール株式会社 | High-tensile-strength hot-rolled plated steel sheet and method for producing same |
JP2013133499A (en) * | 2011-12-27 | 2013-07-08 | Jfe Steel Corp | High-tensile hot-rolled steel sheet excellent in bendability and manufacturing method therefor |
JP2013133525A (en) * | 2011-12-27 | 2013-07-08 | Jfe Steel Corp | High-tensile hot-rolled steel sheet excellent in blanking property and stretch flange formability, and manufacturing method therefor |
Also Published As
Publication number | Publication date |
---|---|
JP6224704B2 (en) | 2017-11-01 |
WO2015118864A1 (en) | 2015-08-13 |
Similar Documents
Publication | Publication Date | Title |
---|---|---|
JP6075517B1 (en) | Hot-rolled steel sheet and manufacturing method thereof | |
JP5765080B2 (en) | High-strength hot-rolled steel sheet excellent in stretch flangeability and manufacturing method thereof | |
JP5423191B2 (en) | High strength steel plate and manufacturing method thereof | |
JP5541263B2 (en) | High-strength hot-rolled steel sheet excellent in workability and manufacturing method thereof | |
JP5321671B2 (en) | High-tensile hot-rolled steel sheet with excellent strength and workability uniformity and method for producing the same | |
JP5884476B2 (en) | High-tensile hot-rolled steel sheet excellent in bending workability and manufacturing method thereof | |
JP5892147B2 (en) | High strength hot rolled steel sheet and method for producing the same | |
JP5915412B2 (en) | High strength hot-rolled steel sheet excellent in bendability and manufacturing method thereof | |
JP5884472B2 (en) | High-strength hot-rolled steel sheet excellent in stretch flangeability and manufacturing method thereof | |
JP5637225B2 (en) | High-strength hot-rolled steel sheet excellent in burring workability and manufacturing method thereof | |
WO2020026594A1 (en) | High-strength hot-rolled plated steel sheet | |
JP6224704B2 (en) | Manufacturing method of high strength hot-rolled steel sheet | |
JP5978614B2 (en) | High-strength hot-rolled steel sheet excellent in punchability and manufacturing method thereof | |
JP5821864B2 (en) | High-strength hot-rolled steel sheet excellent in burring workability and manufacturing method thereof | |
JP5482779B2 (en) | High-tensile hot-rolled steel sheet excellent in punchability and stretch flangeability and manufacturing method thereof | |
JP6610788B2 (en) | High strength hot rolled galvanized steel sheet | |
JP6086080B2 (en) | High-strength cold-rolled steel sheet and manufacturing method thereof | |
JP5594438B2 (en) | High tensile hot rolled galvanized steel sheet and method for producing the same | |
JP2015147957A (en) | High-strength thin steel sheet excellent in shear-plane delayed-fracture resistance characteristic, and production method thereof | |
JP5861434B2 (en) | High-strength hot-rolled steel sheet excellent in punchability and manufacturing method thereof | |
JP2015028199A (en) | High strength hot rolled steel sheet having excellent fatigue resistance and shape freezing properties after slit processing, and manufacturing method thereof | |
JP5884164B2 (en) | High-tensile hot-rolled steel sheet excellent in punchability and stretch flangeability and manufacturing method thereof | |
JP6086078B2 (en) | High-strength cold-rolled steel sheet with excellent stretch flangeability and manufacturing method thereof | |
JP6086077B2 (en) | High-strength cold-rolled steel sheet with excellent workability and method for producing the same | |
JP2013127098A (en) | High-strength hot-dip galvanized steel sheet excellent in workability and method for manufacturing the same |
Legal Events
Date | Code | Title | Description |
---|---|---|---|
A02 | Decision of refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A02 Effective date: 20160329 |
|
A521 | Request for written amendment filed |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20170725 |
|
A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20171005 |
|
R150 | Certificate of patent or registration of utility model |
Ref document number: 6224704 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |
|
R250 | Receipt of annual fees |
Free format text: JAPANESE INTERMEDIATE CODE: R250 |