JP4696853B2 - Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet - Google Patents
Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet Download PDFInfo
- Publication number
- JP4696853B2 JP4696853B2 JP2005316564A JP2005316564A JP4696853B2 JP 4696853 B2 JP4696853 B2 JP 4696853B2 JP 2005316564 A JP2005316564 A JP 2005316564A JP 2005316564 A JP2005316564 A JP 2005316564A JP 4696853 B2 JP4696853 B2 JP 4696853B2
- Authority
- JP
- Japan
- Prior art keywords
- less
- temperature
- steel sheet
- annealing
- ferrite
- 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.)
- Expired - Fee Related
Links
- 229910052799 carbon Inorganic materials 0.000 title claims description 29
- 239000010960 cold rolled steel Substances 0.000 title claims description 29
- 238000004519 manufacturing process Methods 0.000 title claims description 28
- 238000000137 annealing Methods 0.000 claims description 80
- 229910000859 α-Fe Inorganic materials 0.000 claims description 78
- 238000001816 cooling Methods 0.000 claims description 53
- 229910000831 Steel Inorganic materials 0.000 claims description 42
- 239000010959 steel Substances 0.000 claims description 42
- 230000009466 transformation Effects 0.000 claims description 27
- 238000005097 cold rolling Methods 0.000 claims description 23
- OKTJSMMVPCPJKN-UHFFFAOYSA-N Carbon Chemical compound [C] OKTJSMMVPCPJKN-UHFFFAOYSA-N 0.000 claims description 22
- 238000000034 method Methods 0.000 claims description 22
- 238000004804 winding Methods 0.000 claims description 15
- 238000002791 soaking Methods 0.000 claims description 8
- 238000005554 pickling Methods 0.000 claims description 7
- XEEYBQQBJWHFJM-UHFFFAOYSA-N Iron Chemical compound [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 claims description 5
- 229910052750 molybdenum Inorganic materials 0.000 claims description 5
- 229910052758 niobium Inorganic materials 0.000 claims description 5
- 229910052719 titanium Inorganic materials 0.000 claims description 5
- 239000012535 impurity Substances 0.000 claims description 4
- 229910052748 manganese Inorganic materials 0.000 claims description 3
- 229910052757 nitrogen Inorganic materials 0.000 claims description 3
- 229910052742 iron Inorganic materials 0.000 claims description 2
- 229910052698 phosphorus Inorganic materials 0.000 claims description 2
- 229910052804 chromium Inorganic materials 0.000 claims 1
- 150000001247 metal acetylides Chemical class 0.000 description 23
- 239000002245 particle Substances 0.000 description 23
- 238000005098 hot rolling Methods 0.000 description 20
- 238000005096 rolling process Methods 0.000 description 13
- 229910000677 High-carbon steel Inorganic materials 0.000 description 9
- 230000000694 effects Effects 0.000 description 9
- 229910001562 pearlite Inorganic materials 0.000 description 9
- 239000006185 dispersion Substances 0.000 description 8
- 238000001953 recrystallisation Methods 0.000 description 8
- 238000010791 quenching Methods 0.000 description 7
- 230000000171 quenching effect Effects 0.000 description 7
- 230000000052 comparative effect Effects 0.000 description 5
- 230000006866 deterioration Effects 0.000 description 5
- 238000010438 heat treatment Methods 0.000 description 5
- 229910001563 bainite Inorganic materials 0.000 description 4
- 239000000463 material Substances 0.000 description 4
- 239000000203 mixture Substances 0.000 description 4
- 238000012545 processing Methods 0.000 description 4
- 238000004080 punching Methods 0.000 description 4
- 238000005728 strengthening Methods 0.000 description 4
- 230000015572 biosynthetic process Effects 0.000 description 3
- 239000005539 carbonized material Substances 0.000 description 3
- 230000007423 decrease Effects 0.000 description 3
- 238000005498 polishing Methods 0.000 description 3
- 238000010583 slow cooling Methods 0.000 description 3
- 239000006104 solid solution Substances 0.000 description 3
- 239000000126 substance Substances 0.000 description 3
- 238000012360 testing method Methods 0.000 description 3
- 241000446313 Lamella Species 0.000 description 2
- 239000000654 additive Substances 0.000 description 2
- 230000000996 additive effect Effects 0.000 description 2
- 229910001566 austenite Inorganic materials 0.000 description 2
- 229910052796 boron Inorganic materials 0.000 description 2
- 238000005336 cracking Methods 0.000 description 2
- 238000005520 cutting process Methods 0.000 description 2
- 238000000465 moulding Methods 0.000 description 2
- 229920006395 saturated elastomer Polymers 0.000 description 2
- 238000009864 tensile test Methods 0.000 description 2
- 229910000975 Carbon steel Inorganic materials 0.000 description 1
- 230000002411 adverse Effects 0.000 description 1
- 229910045601 alloy Inorganic materials 0.000 description 1
- 239000000956 alloy Substances 0.000 description 1
- 229910052787 antimony Inorganic materials 0.000 description 1
- 230000005540 biological transmission Effects 0.000 description 1
- 239000010962 carbon steel Substances 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 238000007796 conventional method Methods 0.000 description 1
- 229910052802 copper Inorganic materials 0.000 description 1
- 230000007797 corrosion Effects 0.000 description 1
- 238000005260 corrosion Methods 0.000 description 1
- 239000013078 crystal Substances 0.000 description 1
- 230000007547 defect Effects 0.000 description 1
- 230000003111 delayed effect Effects 0.000 description 1
- 238000011156 evaluation Methods 0.000 description 1
- 239000010419 fine particle Substances 0.000 description 1
- 238000005242 forging Methods 0.000 description 1
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
- 229910000734 martensite Inorganic materials 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 229910052759 nickel Inorganic materials 0.000 description 1
- 239000002244 precipitate Substances 0.000 description 1
- 238000001556 precipitation Methods 0.000 description 1
- 238000004881 precipitation hardening Methods 0.000 description 1
- 238000003672 processing method Methods 0.000 description 1
- 238000011160 research Methods 0.000 description 1
- 239000002436 steel type Substances 0.000 description 1
- 238000004781 supercooling Methods 0.000 description 1
- 238000005496 tempering Methods 0.000 description 1
- 229910052718 tin Inorganic materials 0.000 description 1
- 239000011573 trace mineral Substances 0.000 description 1
- 235000013619 trace mineral Nutrition 0.000 description 1
- 229910052721 tungsten Inorganic materials 0.000 description 1
- 229910052720 vanadium Inorganic materials 0.000 description 1
- 229910052726 zirconium Inorganic materials 0.000 description 1
Landscapes
- Heat Treatment Of Sheet Steel (AREA)
Description
本発明は、連続焼鈍法による加工性に優れた高炭素冷延鋼板の製造方法および高炭素冷延鋼板に関する。 The present invention relates to a method for producing a high carbon cold-rolled steel sheet excellent in workability by a continuous annealing method and a high carbon cold-rolled steel sheet.
高炭素鋼板は自動車駆動系部品をはじめ各種機械構造用部品や軸受け部品、工具等まで幅広く使用される。そして、これらの部品の多くは最終的に焼入れ、焼戻しの熱処理が施され、強度、靭性等の特性が付与される。一方で、種々の複雑な形状に加工されるため、優れた加工性が必要となる。特に、近年では部品製造コスト低減を目的とした工程の簡略化や加工方法の変更にともない素材に対する加工性の要求はますます厳しくなっている。 High carbon steel plates are widely used for various parts such as automobile drive system parts, various machine structural parts, bearing parts, and tools. And most of these parts are finally subjected to heat treatment such as quenching and tempering to impart properties such as strength and toughness. On the other hand, since it is processed into various complicated shapes, excellent workability is required. In particular, in recent years, with the simplification of the process and the change of the processing method for the purpose of reducing the part manufacturing cost, the demand for workability of the material has become increasingly severe.
例えば、従来、鍛造、鋳造により製造されてきたリングギヤ、ドライブプレート等の自動車部品については、高炭素鋼板を用いてプレスでの一体成形後熱処理する方法が開発され、一部の部品については実用化されている。 For example, for automotive parts such as ring gears and drive plates that have been manufactured by forging and casting, a method of heat treatment after integrated molding with a press using a high carbon steel sheet has been developed, and some parts have been put to practical use. Has been.
難成形性の高炭素鋼板をこれらの部品に適用するためには素材が軟らかく、プレス成形性に優れることが求められる。また、これらの部品の多くは打抜き加工が施されるため、打抜き工具の長寿命化の要求も厳しくなっており、軟質かつ高加工性を有する鋼板が求められている。このような軟質化と加工性を付与するため、通常は、熱間圧延後の鋼板に球状化焼鈍が施される。球状化処理はフェライト中に球状化したセメンタイトを分散させる処理であり、一般的に、箱焼鈍により行われる。また、高炭素冷延鋼板の製造においては、球状化焼鈍だけでなく、冷間圧延後の再結晶焼鈍に対しても箱焼鈍が行われる。箱焼鈍は炉内にコイルを装入した後、点火、昇温、均熱保持(20〜40hr)、徐冷、脱炉の手順で行われるため、多大なエネルギーと時間を要する。そのため、従来から簡略化が望まれている。 In order to apply a difficult-to-form high carbon steel sheet to these parts, the material is required to be soft and excellent in press formability. In addition, since many of these parts are punched, there is a strict requirement for extending the life of the punching tool, and a steel sheet that is soft and has high workability is required. In order to impart such softening and workability, the steel sheet after hot rolling is usually subjected to spheroidizing annealing. The spheroidizing treatment is a treatment for dispersing spheroidized cementite in ferrite, and is generally performed by box annealing. Further, in the production of a high carbon cold-rolled steel sheet, box annealing is performed not only for spheroidizing annealing but also for recrystallization annealing after cold rolling. Since box annealing is performed in the order of ignition, temperature rise, soaking (20-40 hours), slow cooling, and de-furnace after inserting the coil in the furnace, it requires a great deal of energy and time. For this reason, simplification is conventionally desired.
このような背景の中で高炭素冷延鋼板の製造において、温度の均一性が良く、焼鈍時間の短縮が可能な連続焼鈍法を適用する検討が進められてきた。例えば、特許文献1には、高炭素鋼にAr3点以上の温度域で熱間圧延を施し、その後30℃/s以上で冷却し、Ms〜500℃、または(Ms+Mf)/2の温度範囲で巻取り、フェライト+ベイナイト組織あるいはベイナイト+焼戻しマルテンサイト組織とし、10%以上の冷間圧延後、再結晶温度以上A1以下に加熱して30秒〜600秒保持する製造方法が提案されている。 In such a background, in the production of a high carbon cold-rolled steel sheet, studies have been made to apply a continuous annealing method that has good temperature uniformity and can shorten the annealing time. For example, in Patent Document 1, hot rolling is performed on a high carbon steel at a temperature range of Ar3 point or higher, then cooled at 30 ° C / s or higher, and a temperature of Ms to 500 ° C or (Ms + Mf) / 2 Winding range, ferrite + bainite structure or bainite + tempered martensite structure, after cold rolling of 10% or more, a manufacturing method of heating to the recrystallization temperature to A1 or less and holding for 30 seconds to 600 seconds is proposed Yes.
また、特許文献２にはC:0.3〜1.0%、Si:0.01〜0.5%、Mn:0.2〜2.0%、Al:0.005〜0.10%、N:0.008%以下、S:0.01%以下、およびB:0.0030%以下の高炭素鋼板の熱延鋼板をAc1点直上の温度に加熱後徐冷してAr1点以下の温度とする球状化焼鈍を施した後、40〜80%の冷間圧延後、連続焼鈍法により(Ac1点-50)℃〜Ac1点の温度で10〜180秒保持する製造方法が提案されている。
しかしながら、特許文献１に記載の技術では、熱延組織が初析フェライト+ベイナイト組織からなるため、球状化焼鈍後は初析フェライトと球状炭化物を含むフェライトが混在した組織となり、両者の変形量が大きく異なる。そのため、引張試験時にこれら変形量が大きく異なる粒の粒界に応力が集中し、球状化組織とフェライトの界面にボイドが発生し、破断するため、延性が劣化すると考えられる。また、伸びフランジ性に関しても、初析フェライトと球状炭化物を含むフェライトでは変形量が大きく異なるため、打抜き加工時に打抜き端面近傍には球状化組織と初析フェライトの界面にボイドが発生し、これがクラックに成長し、伸びフランジ性を劣化させると考えられる。また、(Ms+Mf)/2の温度範囲で巻取った場合、球状化焼鈍後のフェライト粒径が細粒となり、強度上昇に起因した加工性の劣化が懸念される。 However, in the technique described in Patent Document 1, since the hot-rolled structure is composed of pro-eutectoid ferrite + bainite structure, after spheroidizing annealing, it becomes a structure in which pro-eutectoid ferrite and ferrite containing spherical carbide are mixed, and the deformation amount of both is reduced. to differ greatly. For this reason, stress is concentrated at the grain boundaries of the grains having greatly different deformation amounts during the tensile test, voids are generated at the interface between the spheroidized structure and the ferrite, and fracture occurs. Also, with regard to stretch flangeability, the amount of deformation differs greatly between proeutectoid ferrite and ferrite containing spherical carbides, so that voids are generated at the interface between the spheroidized structure and proeutectoid ferrite in the vicinity of the punching end surface during punching. It is thought that it grows to deteriorate the stretch flangeability. In addition, when the coil is wound in the temperature range of (Ms + Mf) / 2, the ferrite grain size after spheroidizing annealing becomes fine, and there is a concern about the deterioration of workability due to the strength increase.
また、特許文献２に記載の技術では、熱延条件に対して明確な記載が無く、実施例から通常の熱間圧延と推定される。そして、この実施例記載の条件では、熱延組織が初析フェライト+粗大パーライト組織となるため、球状化焼鈍後は初析フェライトと球状炭化物を含むフェライトが混在した組織となり、延性が劣化すると考えられる。また、粗大パーライト組織を有する場合、未球状炭化物が残存した組織となりやすい。この対策としては焼鈍時間の長時間化が考えられるが、コストが増大するという問題がある。さらに実施例では、完全球状化のため、熱延板冷間圧延を行っているが、これも製造工程が長くなり、コストが増大する。 Moreover, in the technique of patent document 2, there is no clear description with respect to hot rolling conditions, and it is estimated from an Example that it is normal hot rolling. And under the conditions described in this example, the hot-rolled structure becomes pro-eutectoid ferrite + coarse pearlite structure, so after spheroidizing annealing, it becomes a structure in which pro-eutectoid ferrite and ferrite containing spherical carbide are mixed, and the ductility is considered to deteriorate. It is done. Moreover, when it has a coarse pearlite structure | tissue, it becomes easy to become a structure | tissue with which the nonspherical carbide | carbonized_material remained. As a countermeasure, it can be considered that the annealing time is extended, but there is a problem that the cost increases. Furthermore, in an Example, although hot-rolled sheet cold rolling is performed for perfect spheroidization, this also becomes a manufacturing process long and cost increases.
本発明は、かかる事情に鑑み、連続焼鈍による加工性に優れた高炭素冷延鋼板の製造方法および加工性に優れた高炭素冷延鋼板を提供することを目的とする。 An object of this invention is to provide the high carbon cold-rolled steel plate excellent in the manufacturing method of the high carbon cold rolled steel plate excellent in workability by continuous annealing, and workability in view of this situation.
本発明は、高炭素鋼板の生産性向上について鋭意研究を進める中でなされた。そして、その過程で、組成や炭化物の形状および量のみならず、炭化物の分散状態も、鋼板の延性および伸びフランジ性に大きな影響を及ぼす因子であることを見出した。さらに、熱延仕上後の冷却条件を制御し、冷間圧延前の焼鈍(以下、1次焼鈍と称す)組織を制御することにより、冷間圧延後の焼鈍(以下、2次焼鈍と称す)を連続焼鈍化できることがわかった。 The present invention was made in the midst of diligent research on improving the productivity of high-carbon steel sheets. And in the process, it discovered that not only a composition and the shape and quantity of a carbide | carbonized_material but the dispersion state of a carbide | carbonized_material was a factor which has a big influence on the ductility and stretch flangeability of a steel plate. Furthermore, by controlling the cooling conditions after hot-rolling finish and controlling the annealing before cold rolling (hereinafter referred to as primary annealing), annealing after cold rolling (hereinafter referred to as secondary annealing) It was found that can be continuously annealed.
さらに、本発明では、この知見に基づき、上記組織を制御するための製造方法を検討し、連続焼鈍による加工性に優れた高炭素冷延鋼板の製造方法を確立した。 Furthermore, in this invention, based on this knowledge, the manufacturing method for controlling the said structure | tissue was examined, and the manufacturing method of the high carbon cold-rolled steel plate excellent in the workability by continuous annealing was established.
本発明は、以上の知見に基づきなされたもので、その要旨は以下のとおりである。
[1]質量%で、C：0.2〜0.7 %、Si：0.01〜0.35%、Mn：0.1〜0.9%、P：0.03%以下、S：0.035%以下、Al：0.08%以下、N：0.01%以下、Cr：0.05〜0.30%を含有し、残部が鉄および不可避不純物からなる鋼を、(Ar3変態点-20℃)以上の仕上温度で熱間圧延し、次いで、120℃/秒超えの冷却速度で500℃以上650℃以下の冷却停止温度まで1次冷却し、次いで、2次冷却により500℃以上650℃以下の温度に保持した後、600℃以下の温度で巻取り、酸洗後、箱型焼鈍法により、600℃以上Ac1変態点以下の温度で球状化焼鈍した後、冷圧率30%以上で冷間圧延を行い、次いで、連続焼鈍法により、(Ac1-50℃)以上Ac1変態点以下の温度で均熱時間100〜600秒の焼鈍を行うことを特徴とする加工性に優れた高炭素冷延鋼板の製造方法。
[2]前記［1］において、前記鋼として、さらに、質量%でB:0.0005〜0.0030%、Mo:0.005〜0.5%、Ti:0.005〜0.05%、Nb:0.005〜0.1%の一種または二種以上を含有することを特徴とする加工性に優れた高炭素冷延鋼板の製造方法。
[3][1]または[2]のいずれかに記載の製造方法により製造された高炭素冷延鋼板であって、前記高炭素冷延鋼板は、フェライト平均粒径が2.0μm以上、炭化物平均粒径が0.10μm以上2.0μm未満、粒内炭化物の体積率が10%以下であるフェライト粒の体積率が50%以上である組織を有することを特徴とする加工性に優れた高炭素冷延鋼板。
The present invention has been made based on the above findings, and the gist thereof is as follows.
[1] By mass%, C: 0.2 to 0.7%, Si: 0.01 to 0.35%, Mn: 0.1 to 0.9%, P: 0.03% or less, S: 0.035% or less, Al: 0.08% or less, N: 0.01% Hereinafter, steel containing Cr: 0.05 to 0.30%, the balance being iron and inevitable impurities, is hot-rolled at a finishing temperature of (Ar3 transformation point -20 ° C) or higher, and then cooled to over 120 ° C / sec. Primary cooling to a cooling stop temperature of 500 ° C. or more and 650 ° C. or less at a speed, and then holding at a temperature of 500 ° C. or more and 650 ° C. or less by secondary cooling, winding at a temperature of 600 ° C. or less, after pickling, After spheroidizing annealing at a temperature of 600 ° C or higher and below Ac1 transformation point by box-type annealing method, cold rolling is performed at a cold pressure ratio of 30% or higher, and then (Ac1-50 ° C) or higher by continuous annealing method. A method for producing a high carbon cold-rolled steel sheet having excellent workability, characterized by annealing at a temperature equal to or lower than the transformation point for a soaking time of 100 to 600 seconds.
[2] In the above [1], as the steel, one or two of B: 0.0005 to 0.0030%, Mo: 0.005 to 0.5%, Ti: 0.005 to 0.05%, and Nb: 0.005 to 0.1% in mass%. The manufacturing method of the high carbon cold-rolled steel plate excellent in workability characterized by containing the above.
[3] A high-carbon cold-rolled steel sheet produced by the production method according to any one of [1] or [2], wherein the high-carbon cold-rolled steel sheet has a ferrite average grain size of 2.0 μm or more, a carbide average High carbon cold rolling with excellent workability, characterized by having a structure in which the grain size is 0.10 μm or more and less than 2.0 μm, the volume fraction of intragranular carbide is 10% or less, and the volume fraction of ferrite grains is 50% or more steel sheet.
なお、本明細書において、鋼の成分を示す％は、すべて質量％である。 In addition, in this specification,% which shows the component of steel is mass% altogether.
本発明は、加工性の向上を図るに当たって、成分組成および製造条件の制御のみならず、フェライト粒径、炭化物粒径、および炭化物の分散状態をも制御することで、プレス成形や穴拡げ加工におけるボイドの発生を抑制し、クラックの成長を遅くすることができる。その結果、加工性に優れた高炭素冷延鋼板が提供可能となる。また、連続焼鈍を活用することで、加工性に優れた高炭素冷延鋼板を効率的に製造することができる。
また、上記高炭素冷延鋼板を用いることにより、ギアに代表される変速機部品等の加工において加工度を高くとることができ、その結果、製造工程を省略して低コストで部品等を製造することが可能となる。
In order to improve workability, the present invention controls not only the component composition and manufacturing conditions but also the ferrite particle size, carbide particle size, and carbide dispersion state in press forming and hole expanding processing. Generation of voids can be suppressed, and crack growth can be slowed. As a result, it is possible to provide a high carbon cold-rolled steel sheet having excellent workability. Moreover, the high carbon cold-rolled steel plate excellent in workability can be efficiently manufactured by utilizing continuous annealing.
In addition, by using the high carbon cold-rolled steel sheet, it is possible to increase the degree of processing in processing of transmission parts such as gears, and as a result, the manufacturing process is omitted and the parts are manufactured at low cost. It becomes possible to do.
本発明は、下記に示す成分組成に制御し、(Ar3変態点-20℃)以上の仕上温度で熱間圧延し、次いで、120℃/秒超えの冷却速度で500℃以上650℃以下の冷却停止温度まで1次冷却し、次いで、2次冷却により500℃以上650℃以下の温度に保持した後、600℃以下の温度で巻取り、酸洗後、箱型焼鈍法により、600℃以上Ac1変態点以下の温度で球状化焼鈍した後、冷圧率30%以上で冷間圧延を行い、次いで、連続焼鈍法により、(Ac1-50℃)以上Ac1変態点以下の温度で均熱時間100〜600秒の焼鈍を行うことにより高炭素冷延鋼板を製造することを特徴とする。そして、このようにして得られた高炭素冷延鋼板は、フェライト平均粒径が2.0μm以上、炭化物平均粒径が0.10μm以上2.0μm未満、粒内炭化物の体積率が10%以下であるフェライト粒の体積率が50%以上である組織を有しており、優れた加工性を有する。さらに、熱延仕上後の冷却条件を制御し、冷間圧延前の焼鈍（以下、1次焼鈍と称す）組織を制御することにより、冷間圧延後の焼鈍（以下、2次焼鈍と称す）の連続焼鈍化が可能となる。 The present invention is controlled to the following component composition, hot-rolled at a finishing temperature of (Ar3 transformation point -20 ° C) or higher, and then cooled to a temperature of 500 ° C or higher and 650 ° C or lower at a cooling rate exceeding 120 ° C / second. Primary cooling to the stop temperature, then holding at a temperature of 500 ° C or higher and 650 ° C or lower by secondary cooling, winding at a temperature of 600 ° C or lower, pickling, and 6001 ° C or higher by box annealing. After spheroidizing annealing at a temperature equal to or lower than the transformation point, cold rolling is performed at a cold pressure ratio of 30% or higher, and then by a continuous annealing method, the soaking time is 100 at a temperature higher than (Ac1-50 ° C) and lower than the Ac1 transformation point. A high carbon cold-rolled steel sheet is produced by annealing for ˜600 seconds. The high carbon cold-rolled steel sheet thus obtained is a ferrite having an average ferrite grain size of 2.0 μm or more, an average carbide grain size of 0.10 μm or more and less than 2.0 μm, and a volume fraction of intragranular carbides of 10% or less. It has a structure in which the volume ratio of grains is 50% or more and has excellent workability. Furthermore, by controlling the cooling conditions after hot rolling and controlling the microstructure before cold rolling (hereinafter referred to as primary annealing), annealing after cold rolling (hereinafter referred to as secondary annealing). Can be continuously annealed.
以下、本発明を詳細に説明する。
まず、本発明における鋼の化学成分の限定理由は以下の通りである。
(1)C：0.2〜0.7%
Cは、炭素鋼において最も基本になる合金元素である。その含有量によって、焼入れ硬さおよび焼鈍状態での炭化物量が大きく変動する。C含有量が0.2%未満では、自動車用部品等に適用する上で十分な焼入れ硬さが得られない。一方、C含有量が0.7%を超えると熱間圧延後の靭性が低下して鋼帯の製造性、ハンドリングが悪くなるとともに、加工度の高い部品への適用が困難となる。したがって、適度な焼入れ硬さと加工性を兼ね備えた鋼板を提供する観点から、C含有量は0.2%以上0.7%以下とする。
Hereinafter, the present invention will be described in detail.
First, the reasons for limiting the chemical components of steel in the present invention are as follows.
(1) C: 0.2-0.7%
C is the most basic alloy element in carbon steel. The quenching hardness and the amount of carbide in the annealed state vary greatly depending on the content. If the C content is less than 0.2%, sufficient quenching hardness cannot be obtained for application to automotive parts and the like. On the other hand, if the C content exceeds 0.7%, the toughness after hot rolling decreases, the steel strip manufacturability and handling deteriorate, and it becomes difficult to apply to parts with high workability. Therefore, from the viewpoint of providing a steel sheet having both appropriate quenching hardness and workability, the C content is set to 0.2% to 0.7%.
(2)Si：0.01〜0.35%
Siは、焼入れ性を向上させる元素である。Siが0.01%未満では焼入れ時の硬さが不足する。一方、Siが0.35%を超えると固溶強化により、フェライトが硬化し、延性および伸びフランジ性が劣化し、成形加工時に割れ発生の原因となる。したがって、適度な焼入れ硬さと加工性を兼ね備えた鋼板を提供する観点から、Si含有量は0.01％以上0.35%以下、好ましくは0.01%以上0.30%以下とする。
(2) Si: 0.01 to 0.35%
Si is an element that improves hardenability. If Si is less than 0.01%, the hardness during quenching is insufficient. On the other hand, if Si exceeds 0.35%, the solid solution strengthens, the ferrite hardens, the ductility and stretch flangeability deteriorate, and causes cracking during molding. Therefore, from the viewpoint of providing a steel sheet having both appropriate quenching hardness and workability, the Si content is set to 0.01% to 0.35%, preferably 0.01% to 0.30%.
(3)Mn：0.1〜0.9%
Mnは、Siと同様に焼入れ性を向上させる元素である。また、SをMnSとして固定し、スラブの熱間割れを防止する重要な元素である。Mnが0.1%未満では、これらの効果が十分に得られず、また焼入れ性は大幅に低下する。一方、Mnが0.9%を超えると固溶強化により、フェライトが硬化し、加工性の劣化を招く。したがって、適度な焼入れ硬さと加工性を兼ね備えた鋼板を提供する観点から、Mn含有量は0.1％以上0.9%以下、好ましくは0.1%以上0.8%以下とする。
(3) Mn: 0.1-0.9%
Mn is an element that improves hardenability like Si. It is an important element that fixes S as MnS and prevents hot cracking of the slab. If Mn is less than 0.1%, these effects cannot be obtained sufficiently, and the hardenability is greatly reduced. On the other hand, if Mn exceeds 0.9%, the ferrite is hardened due to solid solution strengthening, resulting in deterioration of workability. Therefore, from the viewpoint of providing a steel sheet having both appropriate quenching hardness and workability, the Mn content is 0.1% or more and 0.9% or less, preferably 0.1% or more and 0.8% or less.
(4)P：0.03%以下
Pは粒界に偏析し、延性や靭性を劣化させるため、P含有量は0.03%以下、好ましくは0.02%以下とする。
(4) P: 0.03% or less
P segregates at the grain boundaries and deteriorates ductility and toughness. Therefore, the P content is 0.03% or less, preferably 0.02% or less.
(5)S：0.035%以下
Sは、MnとMnSを形成し、延性および伸びフランジ性を劣化させるため、低減しなければならない元素であり、少ない方が好ましい。しかし、S含有量が0.035%までは許容できるため、S含有量は0.035%以下、好ましくは0.030%以下とする。
(5) S: 0.035% or less
Since S forms Mn and MnS and deteriorates ductility and stretch flangeability, it is an element that must be reduced, and is preferably as small as possible. However, since the S content is acceptable up to 0.035%, the S content is 0.035% or less, preferably 0.030% or less.
(6)Al：0.08%以下
Alは過剰に添加するとAlNが多量に析出し、焼入性を低下させるため、Al含有量は0.08%以下とする。
(6) Al: 0.08% or less
If Al is added excessively, a large amount of AlN precipitates and lowers the hardenability, so the Al content should be 0.08% or less.
(7) N：0.01%以下
Nは過剰に含有している場合は延性の低下をもたらすため、N含有量は0.01%以下とする。
(7) N: 0.01% or less
When N is excessively contained, ductility is lowered, so the N content is 0.01% or less.
(8)Cr：0.05〜0.30%
Crは熱間圧延後の冷却中の初析フェライトの生成を抑制し、延性および伸びフランジ性を向上させ、かつ焼入れ性を向上させる重要な元素である。しかし、Cr含有量が0.05%未満では十分な効果が得られない。一方、0.30%を超えて含有しても、焼入れ性は向上するが、初析フェライト生成の抑制効果が飽和するとともに、コスト増となる。したがって、Cr含有量は0.05％以上0.30%以下とする。
(8) Cr: 0.05-0.30%
Cr is an important element that suppresses the formation of pro-eutectoid ferrite during cooling after hot rolling, improves ductility and stretch flangeability, and improves hardenability. However, if the Cr content is less than 0.05%, a sufficient effect cannot be obtained. On the other hand, if the content exceeds 0.30%, the hardenability is improved, but the effect of suppressing the formation of pro-eutectoid ferrite is saturated and the cost is increased. Therefore, the Cr content is 0.05% or more and 0.30% or less.
本発明鋼は、上記の必須添加元素で目的とする特性が得られるが、上記の必須添加元素に加えて、熱延冷却時の初析フェライト生成の抑制、焼入れ性の向上のためB、Mo、Ti、Nbを必要に応じて１種または２種以上で添加してもよい。その場合、それぞれの添加量が0.0005％未満、0.005%未満、0.005%未満、0.005%未満では添加の効果が十分に得られない。一方、B、Mo、Ti、Nbが、それぞれ0.0030％、0.5%、0.05%、0.1%を超えると、効果が飽和し、コスト増となり、さらに固溶強化、析出強化等により強度上昇が大きくなるため、加工性が劣化する。したがって、これらの元素を添加する場合は、Bは0.0005%以上0.0030%以下、Moは0.005%以上0.5%以下、Tiは0.005%以上0.05%以下、Nbは0.005%以上0.1%以下とする。 The steel of the present invention can achieve the desired properties with the above-mentioned essential additive elements, but in addition to the above-mentioned essential additive elements, B, Mo for suppressing the formation of pro-eutectoid ferrite during hot rolling cooling and improving hardenability. Ti, Nb may be added alone or in combination as required. In that case, if the addition amount is less than 0.0005%, less than 0.005%, less than 0.005%, and less than 0.005%, the effect of addition cannot be sufficiently obtained. On the other hand, if B, Mo, Ti, and Nb exceed 0.0030%, 0.5%, 0.05%, and 0.1%, respectively, the effect will be saturated and the cost will increase, and the strength will increase due to solid solution strengthening and precipitation strengthening. Therefore, workability deteriorates. Therefore, when these elements are added, B is 0.0005% to 0.0030%, Mo is 0.005% to 0.5%, Ti is 0.005% to 0.05%, and Nb is 0.005% to 0.1%.
なお、上記以外の残部はFe及び不可避不純物からなる。不可避不純物として、例えば、Oは非金属介在物を形成し品質に悪影響を及ぼすため、0.003％以下に低減するのが望ましい。また、本発明では、本発明の作用効果を害さない微量元素として、Cu、Ni、W、V、Zr、Sn、Sbを0.1%以下の範囲で含有してもよい。 The remainder other than the above consists of Fe and inevitable impurities. As an unavoidable impurity, for example, O forms non-metallic inclusions and adversely affects quality, so it is desirable to reduce it to 0.003% or less. In the present invention, Cu, Ni, W, V, Zr, Sn, and Sb may be contained in a range of 0.1% or less as trace elements that do not impair the effects of the present invention.
次に、本発明の鋼板の組織について説明する。
(1)フェライト平均粒径：2.0μm以上
フェライト平均粒径（フェライト粒の平均粒径）が2.0μm未満の微細粒となると強度上昇が顕著となり、プレス加工時の負荷が増大する。また、強度の上昇にともない、延性の低下を招き、加工性が劣化する。以上の理由により、フェライト平均粒径2.0μm以上とする。一方、フェライト平均粒径の上限は特に規定しないが、10μm超えでは、打抜き端面性状およびプレス加工後の表面性状が劣化するため、10μm以下とすることが好ましい。なお、フェライト平均粒径は、後述のように製造条件、特に熱間圧延後の1次冷却停止温度、2次冷却保持温度および巻取温度により、制御することができる。
Next, the structure of the steel sheet of the present invention will be described.
(1) Average ferrite particle diameter: 2.0 μm or more When the ferrite average particle diameter (average particle diameter of ferrite grains) is a fine particle of less than 2.0 μm, the strength rises significantly, and the load during press working increases. Further, as the strength increases, ductility is reduced and workability deteriorates. For the above reasons, the ferrite average particle size is 2.0 μm or more. On the other hand, the upper limit of the average ferrite grain size is not particularly specified, but if it exceeds 10 μm, the punched end face properties and the surface properties after press working are deteriorated, so it is preferably 10 μm or less. The ferrite average particle diameter can be controlled by the production conditions, particularly the primary cooling stop temperature after the hot rolling, the secondary cooling holding temperature, and the winding temperature, as will be described later.
(2)炭化物平均粒径：0.10μm以上2.0μm未満
炭化物平均粒径は、加工性一般および穴拡げ加工におけるボイドの発生に大きく影響するため、重要な要素である。炭化物が微細になるとボイドの発生は抑制できるが、炭化物平均粒径が0.10μm未満になると、硬さの上昇に伴い延性が低下し、伸びフランジ性も劣化する。一方、炭化物平均粒径の増加にともない加工性一般は向上するが、2.0μm以上になると、穴拡げ加工におけるボイドの発生により伸びフランジ性が劣化する。以上より、炭化物平均粒径は0.10μm以上2.0μm未満とする。なお、炭化物平均粒径は、後述のように製造条件、特に熱間圧延後の1次冷却停止温度、2次冷却保持温度、巻取温度、そして焼鈍条件により、制御することができる。
(2) Carbide average particle size: 0.10 μm or more and less than 2.0 μm The carbide average particle size is an important factor because it greatly affects the workability in general and the generation of voids in hole expansion processing. Generation of voids can be suppressed when the carbide becomes fine, but when the average particle size of the carbide is less than 0.10 μm, the ductility decreases as the hardness increases, and the stretch flangeability also deteriorates. On the other hand, the workability in general improves with an increase in the average carbide particle diameter, but when it exceeds 2.0 μm, the stretch flangeability deteriorates due to the generation of voids in the hole expanding process. From the above, the carbide average particle size is set to 0.10 μm or more and less than 2.0 μm. The carbide average particle size can be controlled by the production conditions, particularly the primary cooling stop temperature after the hot rolling, the secondary cooling holding temperature, the coiling temperature, and the annealing conditions as described later.
(3)粒内炭化物の体積率が10%以下であるフェライト粒の体積率が50%以上
炭化物の分散状態は、加工性一般および穴拡げ加工におけるボイドの発生に大きく影響するため、重要な要素である。粒内炭化物の体積率が10%超えであるフェライト粒、すなわち、フェライト粒内に炭化物が微細分散したフェライト粒は、引張変形および穴拡げ加工時にフェライトと炭化物の界面にボイドが発生しやすく、また、炭化物の粒子間距離が短いために、発生したボイドが連結しやすい。さらに、炭化物が微細分散することで、硬さの上昇が著しく、延性や伸びフランジ性が劣位となる。一方、粒内炭化物の体積率が10%以下であるフェライト粒の体積率を50%以上とすることで、鋼板硬さが低下し、さらに、ボイドの発生およびボイドの連結が抑制され、延性および伸びフランジ性が大幅に向上する。よって、本発明では、粒内炭化物の体積率が10%以下であるフェライト粒の体積率を50%以上とする。
(3) The volume fraction of the intragranular carbide is 10% or less The volume fraction of the ferrite grain is 50% or more The dispersion state of the carbide greatly affects the workability in general and the generation of voids in the hole expanding process. It is. Ferrite grains with a volume fraction of intragranular carbide exceeding 10%, i.e. ferrite grains with finely dispersed carbides in ferrite grains, tend to generate voids at the interface between ferrite and carbide during tensile deformation and hole expansion. Since the distance between carbide particles is short, the generated voids are easily connected. Furthermore, when the carbide is finely dispersed, the hardness is remarkably increased, and the ductility and stretch flangeability are inferior. On the other hand, by setting the volume fraction of ferrite grains whose volume fraction of intragranular carbide is 10% or less to 50% or more, the steel sheet hardness is reduced, and further, generation of voids and connection of voids are suppressed, ductility and Stretch flangeability is greatly improved. Therefore, in the present invention, the volume fraction of ferrite grains in which the volume fraction of intragranular carbides is 10% or less is set to 50% or more.
ここで、上記理由からフェライト粒内には炭化物を含まないことが好ましく、粒内炭化物を含まないフェライト粒の体積率を50％以上とすることが最も好ましい形態ではある。しかし、粒内炭化物の体積率が10%以下のフェライト粒は、ボイドの発生および連結を抑制する効果が十分に得られ、かつ、炭化物の分散強化による硬さの上昇もなく、延性や伸びフランジ性が良好であり、粒内炭化物を含まないフェライト粒と実質的に同じとみなせる。よって、上記のように、粒内炭化物の体積率が10%以下であるフェライト粒の体積率を50%以上とした。なお、炭化物の分散状態、すなわち、粒内炭化物の体積率が10%以下のフェライト粒の体積率は、製造条件、特に熱間圧延後の1次冷却停止温度、2次冷却保持温度、巻取温度、および焼鈍温度により制御することができる。 Here, for the reasons described above, it is preferable that the ferrite grains do not contain carbides, and the most preferable form is that the volume fraction of ferrite grains not containing intragranular carbides is 50% or more. However, ferrite grains with a volume fraction of intragranular carbide of 10% or less are sufficiently effective in suppressing the generation and connection of voids, and there is no increase in hardness due to dispersion strengthening of carbides. Therefore, it can be regarded as substantially the same as ferrite grains containing no intra-grain carbide. Therefore, as described above, the volume fraction of ferrite grains in which the volume fraction of intragranular carbides is 10% or less was set to 50% or more. Note that the dispersion state of the carbide, that is, the volume fraction of the ferrite grains having a volume fraction of intragranular carbide of 10% or less is determined according to the manufacturing conditions, particularly the primary cooling stop temperature, the secondary cooling holding temperature, and the winding after hot rolling. It can be controlled by temperature and annealing temperature.
なお、フェライト粒内の炭化物の体積率は、鋼板試料の板厚断面を研磨し、ナイタールで腐食後、走査電子顕微鏡で2000倍で、約3000個のフェライト粒を観察し、各フェライト粒についてのフェライトと粒内炭化物の面積比を求め、それを体積率とみなすことにより求めることができる。 The volume fraction of carbides in the ferrite grains was determined by polishing the plate thickness section of the steel sheet sample, corroding with nital, and observing about 3000 ferrite grains with a scanning electron microscope at a magnification of 2000 times. It can be obtained by determining the area ratio of ferrite and intragranular carbide and considering it as the volume fraction.
次に、本発明の加工性に優れた高炭素冷延鋼板の製造方法について説明する。 Next, the manufacturing method of the high carbon cold-rolled steel plate excellent in workability of this invention is demonstrated.
本発明の高炭素冷延鋼板は、上記化学成分範囲に調整された鋼を、(Ar3変態点-20℃)以上の仕上温度で熱間圧延し、次いで、120℃/秒超えの冷却速度で500℃以上650℃以下の冷却停止温度まで1次冷却し、次いで、2次冷却により500℃以上650℃以下の温度に保持した後、600℃以下の温度で巻取り、酸洗後、箱型焼鈍法により、600℃以上Ac1変態点以下の温度で球状化焼鈍した後、冷圧率30%以上で冷間圧延を行い、次いで、連続焼鈍法により、(Ac1-50℃)以上Ac1変態点以下の温度で均熱時間100〜600秒の焼鈍を行うことにより得られる。これについて以下に詳細に説明する。 The high carbon cold-rolled steel sheet of the present invention is obtained by hot rolling a steel adjusted to the above chemical composition range at a finishing temperature of (Ar3 transformation point -20 ° C) or higher, and then at a cooling rate exceeding 120 ° C / second. Primary cooling to a cooling stop temperature of 500 ° C or higher and 650 ° C or lower, then holding at a temperature of 500 ° C or higher and 650 ° C or lower by secondary cooling, winding at a temperature of 600 ° C or lower, pickling, and box-type After spheroidizing annealing at a temperature of 600 ° C or higher and lower than Ac1 transformation point by annealing method, cold rolling is performed at a cold pressure ratio of 30% or higher, and then (Ac1-50 ° C) or higher by the continuous annealing method. It can be obtained by annealing at the following temperature for a soaking time of 100 to 600 seconds. This will be described in detail below.
(1)仕上温度：(Ar3変態点-20℃)以上
鋼を熱間圧延する際の仕上温度が(Ar3変態点-20℃)未満では、一部でフェライト変態が進行するため、初析フェライト粒が増加し、初析フェライトと球状炭化物を含むフェライトの界面にボイドが発生しやすく、延性および伸びフランジ性が劣化する。そこで、(Ar3変態点-20℃)以上の仕上温度で仕上圧延する。これにより、組織の均一化を図ることができ、延性や伸びフランジ性の劣化を抑制できる。仕上温度の上限は特に規定しないが、1000℃を超えるような高温の場合、スケール性欠陥が発生し易くなるため、1000℃以下が好ましい。なお、Ar3変態点（℃）は次の式で算出することができる。
Ar3=930.21-394.75C+54.99Si-14.40Mn+5.77Cr (1)
ここで、式中の元素記号はそれぞれの元素の含有量（質量％）を表す。
なお、圧延負荷の観点からは、仕上温度は高いほうがよく、700℃以上とすることが好ましく、750℃以上とすることがさらに好ましい。
(1) Finishing temperature: (Ar3 transformation point -20 ° C) or higher If the finishing temperature when hot rolling the steel is less than (Ar 3 transformation point -20 ° C), ferrite transformation proceeds in part, and therefore Ferrite grains increase, voids are likely to occur at the interface between pro-eutectoid ferrite and ferrite containing spherical carbides, and ductility and stretch flangeability deteriorate. Therefore, finish rolling is performed at a finishing temperature of (Ar 3 transformation point −20 ° C.) or higher. Thereby, a structure | tissue can be equalized and deterioration of ductility and stretch flangeability can be suppressed. The upper limit of the finishing temperature is not particularly specified, but at a high temperature exceeding 1000 ° C., a scale defect is likely to occur. The Ar 3 transformation point (° C.) can be calculated by the following formula.
Ar 3 = 930.21-394.75C + 54.99Si-14.40Mn + 5.77Cr (1)
Here, the element symbol in a formula represents content (mass%) of each element.
From the viewpoint of rolling load, the finishing temperature should be higher, preferably 700 ° C. or higher, and more preferably 750 ° C. or higher.
(2)1次冷却速度：120℃/秒超え
熱間圧延後の1次冷却方法が徐冷であると、オーステナイトの過冷度が小さく初析フェライトが多く生成する。冷却速度が120℃/秒以下の場合、初析フェライトの生成が顕著となり、初析フェライトと球状炭化物を含むフェライトの界面にボイドが発生しやすく、延性および伸びフランジ性が劣化する。また、パーライトのコロニーおよびラメラ間隔が増大し、球状化焼鈍時間の長時間化を招き、コストが増大する。従って、熱間圧延後の冷却の冷却速度は120℃/秒超とする。なお、冷却速度の上限は特に制限しないが、例えば、現状の設備上の能力からは700℃/秒である。
ここで、冷却速度とは仕上圧延後の冷却開始から冷却停止までの平均冷却速度である。また、仕上圧延後、0.1秒を超え1.0秒未満の時間内で冷却を開始することは、変態後のフェライト結晶粒を微細化し、加工性をより一層向上する上で好ましい。
(2) Primary cooling rate: over 120 ° C / sec. If the primary cooling method after hot rolling is slow cooling, the degree of supercooling of austenite is small and a large amount of proeutectoid ferrite is generated. When the cooling rate is 120 ° C./second or less, pro-eutectoid ferrite is prominent, voids are likely to occur at the interface between pro-eutectoid ferrite and ferrite containing spherical carbides, and ductility and stretch flangeability deteriorate. Further, the spacing between pearlite colonies and lamellae is increased, resulting in longer spheroidizing annealing time and increased cost. Therefore, the cooling rate for cooling after hot rolling is set to exceed 120 ° C./second. The upper limit of the cooling rate is not particularly limited, but for example, it is 700 ° C./second from the current facility capacity.
Here, the cooling rate is an average cooling rate from the start of cooling after finish rolling to the stop of cooling. In addition, it is preferable to start cooling within a time exceeding 0.1 seconds and less than 1.0 seconds after finish rolling in order to refine the ferrite crystal grains after transformation and further improve workability.
(3)1次冷却停止温度：500℃以上650℃以下
熱間圧延後の1次冷却停止温度が650℃超えの場合、その後の冷却中にフェライトが生成しやすく、また、パーライトのコロニーおよびラメラ間隔が増大し、1次焼鈍後に未球状炭化物が残存しやすい。この未球状炭化物は冷間圧延時に砕かれ、次に行われる2次焼鈍時のフェライトの再結晶を抑制し、微細粒となる。また、未球状炭化物の多くは粒内炭化物となる傾向にある。そして、このような細粒硬化および炭化物の微細分散硬化による強度上昇にともない、延性の低下を招き、加工性が劣化する。したがって、熱間圧延後の1次冷却停止温度は650℃以下とする。一方、1次冷却停止温度が500℃未満では、鋼板の形状が劣化し、また、等軸フェライト粒が得られず、加工性が劣化することがある。よって、1次冷却停止温度は500℃以上とする。
(3) Primary cooling stop temperature: 500 ° C or more and 650 ° C or less If the primary cooling stop temperature after hot rolling exceeds 650 ° C, ferrite is likely to form during the subsequent cooling, and pearlite colonies and lamellae The spacing increases, and non-spherical carbides tend to remain after the primary annealing. This non-spherical carbide is crushed during cold rolling, and suppresses recrystallization of ferrite during the subsequent secondary annealing to become fine grains. Further, most of the non-spherical carbides tend to be intragranular carbides. As the strength is increased by such fine grain hardening and fine dispersion hardening of carbide, the ductility is lowered and workability is deteriorated. Therefore, the primary cooling stop temperature after hot rolling is set to 650 ° C. or less. On the other hand, when the primary cooling stop temperature is less than 500 ° C., the shape of the steel sheet is deteriorated, and equiaxed ferrite grains cannot be obtained, so that workability may be deteriorated. Therefore, the primary cooling stop temperature is set to 500 ° C. or higher.
(4)2次冷却保持温度：500℃以上650℃以下
高炭素鋼板の場合、1次冷却停止後に、初析フェライト変態、パーライト変態、ベイナイト変態に伴い、鋼板温度が上昇することがあり、1次冷却停止温度が650℃以下であっても、1次冷却終了から、巻取までに温度が上昇した場合、初析フェライトが生成するとともに、パーライトのラメラ間隔が粗大化する。そのため、1次焼鈍後の未球状炭化物により、加工性が劣化する。また、1次冷却終了から、巻取までに、温度が500℃未満になると、鋼板の形状が劣化し、また、等軸フェライト粒が得られず、加工性が劣化することがある。これらの理由により、2次冷却により、1次冷却終了から巻取までの温度を制御することは重要であり、1次冷却終了から巻取まで500℃以上650℃以下の温度で保持することとする。このように500℃以上650℃以下の温度で保持することにより、加工性の劣化を防止することができる。なお、この場合の2次冷却はラミナー冷却等により行うことができる。
(4) Secondary cooling holding temperature: In the case of a high carbon steel sheet of 500 ° C or higher and 650 ° C or lower, the steel plate temperature may increase with the proeutectoid ferrite transformation, pearlite transformation, and bainite transformation after the primary cooling is stopped. Even when the secondary cooling stop temperature is 650 ° C. or lower, if the temperature rises from the end of the primary cooling to the winding, proeutectoid ferrite is generated and the pearlite lamella spacing becomes coarse. Therefore, workability deteriorates due to the non-spherical carbide after the primary annealing. In addition, when the temperature is lower than 500 ° C. from the end of the primary cooling to the winding, the shape of the steel sheet is deteriorated, and equiaxed ferrite grains cannot be obtained, so that workability may be deteriorated. For these reasons, it is important to control the temperature from the end of the primary cooling to the winding by the secondary cooling, and keep the temperature from 500 ° C to 650 ° C from the end of the primary cooling to the winding. To do. By maintaining the temperature at 500 ° C. or higher and 650 ° C. or lower as described above, deterioration of workability can be prevented. In this case, the secondary cooling can be performed by laminar cooling or the like.
(5)巻取温度：600℃以下
巻取温度が高いほど、パーライトのラメラ間隔が大きくなる。そのため、巻取温度が600℃超えでは、未球状炭化物が残存しやすく、これにより加工性が劣化することがある。したがって、巻取温度は600℃以下とする。なお、巻取温度の下限は特に規定しないが、低温になるほど鋼板の形状が劣化するため、200℃以上とすることが好ましい。
(5) Winding temperature: 600 ° C. or less The higher the winding temperature, the larger the pearlite lamella spacing. For this reason, when the coiling temperature exceeds 600 ° C., non-spherical carbides are likely to remain, which may deteriorate workability. Therefore, the coiling temperature is 600 ° C. or less. Although the lower limit of the coiling temperature is not particularly defined, the shape of the steel sheet is deteriorated as the temperature is lowered, and is preferably set to 200 ° C. or higher.
(6)酸洗：実施
巻取後の熱延鋼板は、冷間圧延を行う前にスケール除去のため、酸洗を施す。酸洗は常法にしたがって行えばよい。
(6) Pickling: The hot-rolled steel sheet after winding is pickled to remove scale before cold rolling. Pickling may be performed according to a conventional method.
(7)1次焼鈍(球状化焼鈍)温度：600℃以上Ac1変態点以下
熱延鋼板を酸洗した後、冷間圧延を行うが、その前に炭化物を球状化するために1次焼鈍(球状化焼鈍)を行う。この1次焼鈍温度が600℃未満の場合、炭化物の球状化が不十分となり、焼鈍効果が得られない。一方、1次焼鈍温度がAc1変態点を超える場合、一部がオーステナイト化し、冷却中に再度パーライトを生成するため、球状化組織が得られない。以上より、1次焼鈍温度は600℃以上Ac1変態点以下とする。なお、本発明では冷間圧延後、2次焼鈍を連続焼鈍法にて行う。そのため、炭化物の球状化率は高いほうが好ましく、80%以上とすることが好ましい。球状化率を80%以上とするためには、1次焼鈍温度を680℃以上とすることが好ましく、1次焼鈍時間を40時間超えとすることが好ましい。ここで、球状化率は、全炭化物数に対するアスペクト比が3以下である球状炭化物数の割合とする。また、本発明では、炭化物の完全球状化のため、1次焼鈍は箱型焼鈍法により行うものとする。
(7) Primary annealing (spheroidizing annealing) temperature: 600 ° C or more and Ac1 transformation point or less After pickling hot rolled steel sheet, cold rolling is performed, but before that, primary annealing is performed to spheroidize carbides ( Spheroidizing annealing). When the primary annealing temperature is less than 600 ° C., the spheroidization of the carbide becomes insufficient and the annealing effect cannot be obtained. On the other hand, when the primary annealing temperature exceeds the Ac 1 transformation point, a part is austenitized and pearlite is generated again during cooling, so that a spheroidized structure cannot be obtained. From the above, the primary annealing temperature is set to 600 ° C. or more and Ac1 transformation point or less. In the present invention, after cold rolling, secondary annealing is performed by a continuous annealing method. Therefore, the spheroidization rate of the carbide is preferably high, and is preferably 80% or more. In order to make the spheroidization rate 80% or more, the primary annealing temperature is preferably 680 ° C. or more, and the primary annealing time is preferably more than 40 hours. Here, the spheroidization rate is the ratio of the number of spherical carbides having an aspect ratio of 3 or less with respect to the total number of carbides. In the present invention, the primary annealing is performed by a box-type annealing method in order to completely spheroidize the carbide.
(8)冷間圧延の圧下率：30%以上
冷間圧延を行うことにより、2次焼鈍時のフェライトの再結晶を助長し、フェライト粒が等軸となり、加工性が向上する。しかし、冷間圧延の圧下率が30％未満では上記効果が得られないばかりか、2次焼鈍後に未再結晶部が残存し、かえって加工性を劣化させる。したがって、冷間圧延の圧下率を30%以上とする。なお、圧下率の上限は特に制約はないが、圧延負荷の問題から80％以下とすることが好ましい。
(8) Cold rolling reduction ratio: By performing cold rolling for 30% or more, recrystallization of ferrite during secondary annealing is promoted, ferrite grains become equiaxed, and workability is improved. However, if the rolling reduction of cold rolling is less than 30%, not only the above effects can be obtained, but also non-recrystallized portions remain after secondary annealing, and workability is deteriorated. Therefore, the rolling reduction of cold rolling is set to 30% or more. The upper limit of the rolling reduction is not particularly limited, but is preferably 80% or less from the viewpoint of rolling load.
(9)2次焼鈍(再結晶焼鈍)：(Ac1-50℃)以上Ac1変態点以下の温度で100〜600秒均熱
冷間圧延後、フェライト再結晶のために2次焼鈍(再結晶焼鈍)を行う。2次焼鈍温度が(Ac1-50℃)未満の場合、または、2次焼鈍温度が本発明範囲内であっても均熱時間が100秒未満の場合、軟化もしくは再結晶が不十分となる。したがって、2次焼鈍温度は(Ac1-50℃)以上、均熱時間は100秒以上、好ましくは180秒以上とする。一方で、2次焼鈍温度がAc1変態点を超える場合、一部がオーステナイト化し、冷却中に再度パーライトが生成するため、硬度が高くなり加工性が劣化する。また、2次焼鈍の目的はフェライトが再結晶さえすればよく、長時間の焼鈍は効果が飽和する。よって、2次焼鈍温度はAc1変態点以下、均熱時間は600秒以下とする。なお、Ac1変態点（℃）は次の式で算出することができる。
Ac1=754.83-32.25C+23.32Si-17.76Mn+17.13Cr (2)
ここで、式中の元素記号はそれぞれの元素の含有量（質量％）を表す。
なお、2次焼鈍は、連続焼鈍法にて行うものとする。
(9) Secondary annealing (recrystallization annealing): Secondary annealing (recrystallization annealing) for ferrite recrystallization after isothermal cold rolling at a temperature of (Ac1-50 ° C) to the Ac1 transformation temperature for 100 to 600 seconds. )I do. When the secondary annealing temperature is less than (Ac1-50 ° C.), or when the soaking time is less than 100 seconds even if the secondary annealing temperature is within the range of the present invention, softening or recrystallization becomes insufficient. Therefore, the secondary annealing temperature is (Ac1-50 ° C.) or more, and the soaking time is 100 seconds or more, preferably 180 seconds or more. On the other hand, when the secondary annealing temperature exceeds the Ac 1 transformation point, part of it becomes austenite and pearlite is generated again during cooling, so that the hardness increases and the workability deteriorates. The purpose of the secondary annealing is only to recrystallize the ferrite, and the effect of saturation for a long time annealing. Therefore, the secondary annealing temperature is set to Ac1 transformation point or less, and the soaking time is set to 600 seconds or less. The Ac 1 transformation point (° C.) can be calculated by the following formula.
Ac 1 = 754.83-32.25C + 23.32Si-17.76Mn + 17.13Cr (2)
Here, the element symbol in a formula represents content (mass%) of each element.
The secondary annealing is performed by a continuous annealing method.
本発明の高炭素鋼の成分調製には、転炉あるいは電気炉のどちらでも使用可能である。このように成分調製された高炭素鋼を、造塊−分塊圧延または連続鋳造により鋼素材である鋼スラブとする。この鋼スラブについて熱間圧延を行うが、その際、スラブ加熱温度は、スケール発生による表面状態の劣化を避けるため1300℃以下とすることが好ましい。 Either a converter or an electric furnace can be used for preparing the components of the high carbon steel of the present invention. The high carbon steel whose components are prepared in this way is used as a steel slab, which is a steel material, by ingot-bundling rolling or continuous casting. The steel slab is hot-rolled, and at that time, the slab heating temperature is preferably 1300 ° C. or lower in order to avoid deterioration of the surface state due to generation of scale.
なお、熱間圧延時に粗圧延を省略して仕上圧延を行ってもよく、連続鋳造スラブをそのまま又は温度低下を抑制する目的で保熱しつつ圧延する直送圧延を行ってもよい。また、仕上温度確保のため、熱間圧延中にバーヒータ等の加熱手段により圧延材の加熱を行ってもよい。なお、球状化促進あるいは硬度低減のため、巻取後にコイルを徐冷カバー等の手段で保温してもよい。 In addition, rough rolling may be omitted during hot rolling and finish rolling may be performed, or direct feed rolling may be performed in which a continuously cast slab is rolled as it is or for the purpose of suppressing temperature reduction. In order to secure the finishing temperature, the rolled material may be heated by a heating means such as a bar heater during hot rolling. In order to promote spheroidization or reduce hardness, the coil may be kept warm by means such as a slow cooling cover after winding.
このようにして得られた高炭素冷延鋼板が、優れた加工性を有する理由は次のように考えられる。延性および伸びフランジ性には、鋼板および打抜き端面の内部組織が大きく影響する。特に、粒内炭化物の体積率が10%超えのフェライト粒が多い(粒内炭化物が分散した組織)場合、炭化物の粒子間距離が小さいため、発生したボイドの連結が速く、クラックの進展が速いことが確認されている。一方、粒内炭化物の体積率が10%以内のフェライト粒を50%以上とすることでクラックの進展が遅延することが確認されている。
このように製造条件の制御のみならず、炭化物平均粒径および炭化物の分散状態を制御することにより、ボイドの連結および成長を抑制することができる。
The reason why the high carbon cold rolled steel sheet thus obtained has excellent workability is considered as follows. Ductility and stretch flangeability are greatly affected by the internal structure of the steel plate and the punched end face. In particular, when there are many ferrite grains with a volume fraction of intragranular carbide exceeding 10% (structure in which intragranular carbides are dispersed), the distance between the carbide particles is small, so that the generated voids are connected quickly and the crack progresses quickly. It has been confirmed. On the other hand, it has been confirmed that the growth of cracks is delayed by setting the ferrite grains having a volume fraction of intragranular carbides within 10% to 50% or more.
Thus, not only the control of the manufacturing conditions, but also the connection and growth of voids can be suppressed by controlling the carbide average particle size and the dispersion state of the carbides.
表1に示す化学成分を有する鋼の連続鋳造スラブを1250℃に加熱し、表2に示す条件にて熱間圧延、箱型焼鈍法による一次焼鈍、冷間圧延、および連続焼鈍法による二次焼鈍を行い、板厚3.0mmの冷延鋼板を製造した。ここで、鋼板No.1〜3は製造条件が本発明範囲内の本発明例であり、鋼板No.4〜11は製造条件が本発明から外れる比較例、鋼板No.12、13は鋼成分が本発明から外れる比較例である。 A continuous cast slab of steel having the chemical components shown in Table 1 is heated to 1250 ° C and hot rolled under the conditions shown in Table 2, primary annealing by box annealing, cold rolling, and secondary by continuous annealing. Annealing was performed to produce a cold-rolled steel sheet having a thickness of 3.0 mm. Here, steel plates No. 1 to 3 are examples of the present invention in which the manufacturing conditions are within the scope of the present invention, steel plates No. 4 to 11 are comparative examples in which the manufacturing conditions deviate from the present invention, and steel plates No. 12 and 13 are steel components. Is a comparative example deviating from the present invention.
次に、上記により得られた冷延鋼板からサンプルを採取し、フェライト平均粒径、炭化物平均粒径ならびに分散状態を測定し、性能評価のため、硬度、伸び(JIS5号C方向)および伸びフランジ性を測定した。それぞれの測定方法および条件については以下の通りである。 Next, a sample was taken from the cold-rolled steel sheet obtained as described above, and the ferrite average particle size, carbide average particle size and dispersion state were measured, and for performance evaluation, hardness, elongation (JIS5C direction) and stretch flange Sex was measured. Each measuring method and conditions are as follows.
＜フェライト平均粒径＞
サンプルの板厚断面での光顕組織から，JIS G 0552に記載の切断法により行った。
<Ferrite average particle size>
It was carried out by the cutting method described in JIS G 0552 from the light microscopic structure in the plate thickness section of the sample.
＜炭化物平均粒径＞
サンプルの板厚断面を研磨・腐食後、走査型電子顕微鏡にてミクロ組織を撮影し、50μm×50μmの範囲で炭化物粒径の測定を行った。
<Carbide average particle size>
After polishing and corrosion of the plate thickness section of the sample, the microstructure was photographed with a scanning electron microscope, and the carbide particle size was measured in the range of 50 μm × 50 μm.
＜炭化物の分散状態(粒内炭化物の体積率が10%以下であるフェライト粒の体積率)＞
サンプルの板厚断面を研磨・腐食後、走査型電子顕微鏡にて約2000倍で、約3000個のフェライト粒を観察し、各フェライト粒について、フェライトの面積と粒内炭化物の面積比により求めた。
<Carbide dispersion state (volume fraction of ferrite grains in which the volume fraction of intra-grain carbide is 10% or less)>
After polishing and corroding the plate thickness section of the sample, about 3,000 ferrite grains were observed with a scanning electron microscope at about 2000 times, and each ferrite grain was determined by the ratio of the area of ferrite and the area of intragranular carbides. .
＜硬度＞
試料の切断面をバフ研磨仕上後、板厚中央部にて荷重500gfの条件下でヴィッカース硬さ(Hv)を測定した。
<Hardness>
The cut surface of the sample was buffed, and then the Vickers hardness (Hv) was measured at the center of the plate thickness under a load of 500 gf.
＜伸び(JIS5号C方向)＞
圧延方向に対して90°の方向にJIS5号試験片を切り出した後、クロスヘッド10mm/min、標点間距離l0=50mmの条件下で引張試験を行い、試験後の標点間距離lを測定して、次式で定義される伸び量：Elを求めた。
El=100×(l-l0)/l0 (1)
＜伸びフランジ性＞
サンプルを、ポンチ径d0=10mm、ダイス径10.46mm（クリアランス20％）の打抜き工具を用いて打抜き後、穴拡げ試験を実施した。穴拡げ試験は、円筒平底ポンチ（50mmφ、5R）にて押し上げる方法で行い、穴縁に板厚貫通クラックが発生した時点での穴径d1を測定して、次式で定義される穴拡げ率：λ(％)を求めた。
λ＝100×(d1-d0)/d0 (2)
以上の測定により得られた結果を表３に示す。なお、伸びフランジ性は穴拡げ率λで評価した。
<Elongation (JIS5 C direction)>
After cutting out a JIS5 test piece in a direction 90 ° to the rolling direction, a tensile test was performed under the conditions of a crosshead of 10 mm / min and a distance between gauge points l 0 = 50 mm. Was measured, and the elongation amount defined by the following formula: El was obtained.
El = 100 × (ll 0 ) / l 0 (1)
<Stretch flangeability>
A sample was punched with a punching tool having a punch diameter d 0 = 10 mm and a die diameter 10.46 mm (clearance 20%), and then a hole expansion test was performed. Hole expansion test was carried out in a manner to push up at a cylindrical flat bottom punch (50 mm [phi], 5R), to measure the hole diameter d 1 at the time when through thickness cracks are formed in the hole edge, hole expansion is defined by the following equation Rate: λ (%) was determined.
λ = 100 × (d 1 -d 0 ) / d 0 (2)
Table 3 shows the results obtained by the above measurement. The stretch flangeability was evaluated by the hole expansion rate λ.
表3において、鋼板No.1〜3は製造条件が本発明範囲であり、フェライトの平均粒径が2.0μm以上、炭化物平均粒径が0.10μm以上2.0μm未満、粒内炭化物の体積率が10%以下であるフェライト粒の体積率が50%以上の本発明例である。いずれの発明例も、比較例No.4〜11と比べて、それぞれ同じ鋼種について、延性(El)、穴拡げ率(λ)ともに向上しており、優れた加工性を有することがわかる。 In Table 3, the production conditions of steel plates No. 1 to 3 are within the scope of the present invention, the average grain size of ferrite is 2.0 μm or more, the average grain size of carbide is 0.10 μm or more and less than 2.0 μm, and the volume fraction of intragranular carbide is 10 This is an example of the present invention in which the volume fraction of ferrite grains having a percentage of 50% or less is 50% or more. In any of the inventive examples, compared to Comparative Examples Nos. 4 to 11, both the ductility (El) and the hole expansion ratio (λ) are improved for the same steel type, and it can be seen that the invention has excellent workability.
一方、鋼板No.4〜11は製造条件が本発明範囲を外れた比較例、鋼板No.12、13は鋼成分が本発明から外れる比較例である。鋼板No.4〜13のいずれも粒内炭化物の体積率が10%以下であるフェライト粒の体積率が50%未満と、本発明の範囲外となっており、そのため、粒内炭化物が増加し、ボイドの連結が生じやすく、加工性が劣っている。
中でも、フェライト平均粒径が2.0μm未満、かつ炭化物平均粒径が下限である0.10μm未満と、本発明の範囲外である鋼板No.10、11は、強度上昇が顕著となり、延性の低下を招き、加工性が劣っている。
On the other hand, steel plates No. 4 to 11 are comparative examples in which the production conditions are out of the scope of the present invention, and steel plate Nos. 12 and 13 are comparative examples in which the steel components are out of the present invention. Steel plates No. 4 to 13 have a volume fraction of ferrite grains in which the volume fraction of intragranular carbides is 10% or less, which is outside the scope of the present invention. , Voids are easily connected and workability is poor.
Among them, the ferrite average particle size is less than 2.0 μm, and the carbide average particle size is less than 0.10 μm which is the lower limit, and the steel plates No. 10 and 11 which are outside the scope of the present invention have a significant increase in strength and a decrease in ductility. Invited, workability is inferior.
本発明の冷延鋼板は、自動車駆動系部品をはじめとする各種機械構造用部品以外にも、優れた加工性（伸びおよび伸びフランジ性）が要求される用途に対しても好適である。 The cold-rolled steel sheet of the present invention is also suitable for applications requiring excellent workability (elongation and stretch flangeability) in addition to various machine structural parts including automobile drive system parts.
Claims (3)
次いで、120℃/秒超えの冷却速度で500℃以上650℃以下の冷却停止温度まで1次冷却し、次いで、2次冷却により500℃以上650℃以下の温度に保持した後、
600℃以下の温度で巻取り、酸洗後、
箱型焼鈍法により、600℃以上Ac1変態点以下の温度で球状化焼鈍した後、
30%以上の冷圧率で冷間圧延を行い、
次いで、連続焼鈍法により、(Ac1-50℃)以上Ac1変態点以下の温度で均熱時間100〜600秒の焼鈍を行うことを特徴とする加工性に優れた高炭素冷延鋼板の製造方法。 In mass%, C: 0.2 to 0.7%, Si: 0.01 to 0.35%, Mn: 0.1 to 0.9%, P: 0.03% or less, S: 0.035% or less, Al: 0.08% or less, N: 0.01% or less, Cr : A steel containing 0.05 to 0.30%, the balance being iron and inevitable impurities, hot-rolled at a finishing temperature of (Ar3 transformation point -20 ° C) or higher,
Next, primary cooling to a cooling stop temperature of 500 ° C. or more and 650 ° C. or less at a cooling rate exceeding 120 ° C./second, and then holding at a temperature of 500 ° C. or more and 650 ° C. or less by secondary cooling,
After winding at a temperature of 600 ° C or less, pickling,
After spheroidizing annealing at a temperature of 600 ° C or more and Ac1 transformation point or less by the box annealing method,
Cold rolling at a cold pressure ratio of 30% or more,
Next, a method for producing a high carbon cold-rolled steel sheet having excellent workability, characterized by performing annealing at a temperature of (Ac1-50 ° C.) or higher and below the Ac1 transformation point by a continuous annealing method for a soaking time of 100 to 600 seconds. .
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2005316564A JP4696853B2 (en) | 2005-10-31 | 2005-10-31 | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2005316564A JP4696853B2 (en) | 2005-10-31 | 2005-10-31 | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2007119883A JP2007119883A (en) | 2007-05-17 |
| JP4696853B2 true JP4696853B2 (en) | 2011-06-08 |
Family
ID=38144026
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2005316564A Expired - Fee Related JP4696853B2 (en) | 2005-10-31 | 2005-10-31 | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP4696853B2 (en) |
Families Citing this family (8)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP5114189B2 (en) * | 2007-12-26 | 2013-01-09 | 株式会社神戸製鋼所 | Cold-working steel, its manufacturing method and cold-worked steel parts |
| DE102009010442A1 (en) * | 2009-02-26 | 2010-09-02 | C.D. Wälzholz GmbH | Micro-alloyed carbon steel as a texture-rolled strip steel, in particular for spring elements |
| JP6068172B2 (en) * | 2013-02-05 | 2017-01-25 | 株式会社神戸製鋼所 | Soft high carbon steel sheet |
| JP6068291B2 (en) * | 2013-08-07 | 2017-01-25 | 株式会社神戸製鋼所 | Soft high carbon steel sheet |
| ES2750615T3 (en) * | 2014-03-07 | 2020-03-26 | Nippon Steel Corp | Medium / high carbon steel sheet and method for manufacturing it |
| CN111748729A (en) * | 2019-03-27 | 2020-10-09 | 宝山钢铁股份有限公司 | Steel sheet for lid manufacture having excellent sealing properties and internal pressure resistance, and method for producing same |
| JP7534618B2 (en) | 2020-09-18 | 2024-08-15 | 日本製鉄株式会社 | High carbon steel parts |
| JP7534617B2 (en) | 2020-09-18 | 2024-08-15 | 日本製鉄株式会社 | High carbon steel parts |
Citations (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH1088237A (en) * | 1996-09-11 | 1998-04-07 | Sumitomo Metal Ind Ltd | Production of cold rolled high carbon steel strip |
| JPH10147816A (en) * | 1996-09-20 | 1998-06-02 | Sumitomo Metal Ind Ltd | Production of hot rolled high carbon steel plate and cold rolled high carbon steel sheet, excellent in formability |
| JPH1136044A (en) * | 1997-07-16 | 1999-02-09 | Nkk Corp | High carbon steel having excellent blanking workability |
| JP2003013144A (en) * | 2001-06-28 | 2003-01-15 | Nkk Corp | Method for manufacturing high-carbon cold-rolled steel sheet superior in stretch flange formability |
| JP2003089846A (en) * | 2001-09-17 | 2003-03-28 | Nkk Corp | High carbon steel sheet for working having reduced plane anisotropy, and production method therefor |
| JP2005097659A (en) * | 2003-09-24 | 2005-04-14 | Nippon Steel Corp | High carbon steel plate having little in-plane anisotropy |
| JP2005290547A (en) * | 2004-03-10 | 2005-10-20 | Jfe Steel Kk | High carbon hot-rolled steel sheet having excellent ductility and stretch-flange formability, and production method therefor |
-
2005
- 2005-10-31 JP JP2005316564A patent/JP4696853B2/en not_active Expired - Fee Related
Patent Citations (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JPH1088237A (en) * | 1996-09-11 | 1998-04-07 | Sumitomo Metal Ind Ltd | Production of cold rolled high carbon steel strip |
| JPH10147816A (en) * | 1996-09-20 | 1998-06-02 | Sumitomo Metal Ind Ltd | Production of hot rolled high carbon steel plate and cold rolled high carbon steel sheet, excellent in formability |
| JPH1136044A (en) * | 1997-07-16 | 1999-02-09 | Nkk Corp | High carbon steel having excellent blanking workability |
| JP2003013144A (en) * | 2001-06-28 | 2003-01-15 | Nkk Corp | Method for manufacturing high-carbon cold-rolled steel sheet superior in stretch flange formability |
| JP2003089846A (en) * | 2001-09-17 | 2003-03-28 | Nkk Corp | High carbon steel sheet for working having reduced plane anisotropy, and production method therefor |
| JP2005097659A (en) * | 2003-09-24 | 2005-04-14 | Nippon Steel Corp | High carbon steel plate having little in-plane anisotropy |
| JP2005290547A (en) * | 2004-03-10 | 2005-10-20 | Jfe Steel Kk | High carbon hot-rolled steel sheet having excellent ductility and stretch-flange formability, and production method therefor |
Also Published As
| Publication number | Publication date |
|---|---|
| JP2007119883A (en) | 2007-05-17 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| JP5292698B2 (en) | Extremely soft high carbon hot-rolled steel sheet and method for producing the same | |
| KR101912512B1 (en) | High-strength cold-rolled steel sheet and method for manufacturing the same | |
| JP5050433B2 (en) | Method for producing extremely soft high carbon hot-rolled steel sheet | |
| JP4650006B2 (en) | High carbon hot-rolled steel sheet excellent in ductility and stretch flangeability and method for producing the same | |
| US20050199322A1 (en) | High carbon hot-rolled steel sheet and method for manufacturing the same | |
| WO2017179372A1 (en) | High strength steel sheet and manufacturing method therefor | |
| JP5126844B2 (en) | Steel sheet for hot pressing, manufacturing method thereof, and manufacturing method of hot pressed steel sheet member | |
| JP4600196B2 (en) | High carbon cold-rolled steel sheet with excellent workability and manufacturing method thereof | |
| KR20170107057A (en) | High-strength cold-rolled steel plate and method for producing same | |
| JP2008069452A (en) | Hot-rolled high-carbon steel sheet and process for production of the same | |
| JP5640931B2 (en) | Medium carbon cold-rolled steel sheet excellent in workability and hardenability and its manufacturing method | |
| JP4696853B2 (en) | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet | |
| WO2013180180A1 (en) | High strength cold-rolled steel plate and manufacturing method therefor | |
| US11401569B2 (en) | High-strength cold-rolled steel sheet and method for manufacturing same | |
| JP5358914B2 (en) | Super soft high carbon hot rolled steel sheet | |
| JPWO2019151017A1 (en) | High-strength cold-rolled steel sheet, high-strength plated steel sheet, and methods for producing them | |
| JP5302840B2 (en) | High-strength cold-rolled steel sheet with an excellent balance between elongation and stretch flangeability | |
| JP3879446B2 (en) | Method for producing high carbon hot-rolled steel sheet with excellent stretch flangeability | |
| JP4380471B2 (en) | High carbon hot rolled steel sheet and manufacturing method thereof | |
| JP4696753B2 (en) | Method for producing high carbon cold-rolled steel sheet excellent in punching workability and high-carbon cold-rolled steel sheet | |
| JP3879447B2 (en) | Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability | |
| JP4380469B2 (en) | High carbon hot rolled steel sheet and manufacturing method thereof | |
| JP4622609B2 (en) | Method for producing soft high workability high carbon hot rolled steel sheet with excellent stretch flangeability | |
| JP4403925B2 (en) | High carbon cold-rolled steel sheet and method for producing the same | |
| JP5860345B2 (en) | High-strength cold-rolled steel sheet with small variation in mechanical properties and method for producing the same |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20080825 |
|
| A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20110113 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20110201 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20110214 |
|
| R150 | Certificate of patent or registration of utility model |
Ref document number: 4696853 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 |
|
| LAPS | Cancellation because of no payment of annual fees |