JP2007031762A - High-carbon cold-rolled steel sheet superior in workability, and manufacturing method therefor - Google Patents
High-carbon cold-rolled steel sheet superior in workability, and manufacturing method therefor Download PDFInfo
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本発明は、加工性に優れた高炭素冷延鋼板およびその製造方法に関する。 The present invention relates to a high carbon cold-rolled steel sheet excellent in workability and a method for producing the same.
工具あるいは自動車部品（ギア、ミッション）等に使用される高炭素鋼板は、打抜き、成形後、焼入れ焼戻し等の熱処理が施される。これらの部品加工を行うユーザの要求の１つとして、複雑形状に成形するための延性の指標である伸び特性と、打抜き後の成形における穴拡げ加工（バーリング）性の向上がある。これらのうち、穴拡げ加工性は、伸びフランジ性で評価されている。そのため、延性と同時に伸びフランジ性の優れた材料が望まれている。 High carbon steel sheets used for tools or automobile parts (gears, missions) and the like are subjected to heat treatment such as quenching and tempering after punching and forming. One of the requirements of users who perform these parts processing is an elongation characteristic that is an index of ductility for forming into a complex shape, and an improvement in hole expansion processing (burring) in forming after punching. Among these, the hole expansion workability is evaluated by stretch flangeability. Therefore, a material having excellent ductility and stretch flangeability is desired.
このような高炭素鋼板の伸びフランジ性の向上については、いくつかの技術が検討されている。例えば、特許文献１及び特許文献２には、冷間圧延を経たプロセスにおいて、伸びフランジ性に優れた中・高炭素鋼板を作る方法が提案されている。この技術は、Ｃ：0.1〜0.8質量％を含有する鋼からなり、金属組織が実質的にフェライト＋パーライト組織であり、必要に応じて初析フェライト面積率がＣ（質量％）により決まる所定の値以上、パーライトラメラ間隔が0.1μm以上の熱延鋼板に、15％以上の冷間圧延を施し、次いで、３段階又は２段階の温度範囲で長時間保持する３段階又は２段階焼鈍を施すというものである。 Several techniques have been studied for improving the stretch flangeability of such a high-carbon steel sheet. For example, Patent Document 1 and Patent Document 2 propose a method for producing a medium / high carbon steel sheet having excellent stretch flangeability in a process after cold rolling. This technique is made of steel containing C: 0.1 to 0.8% by mass, the metal structure is substantially a ferrite + pearlite structure, and the pro-eutectoid ferrite area ratio is determined by C (% by mass) as required. More than the above, hot rolled steel sheet with a pearlite lamella spacing of 0.1 μm or more is subjected to cold rolling of 15% or more, and then subjected to three-stage or two-stage annealing for a long time in a three-stage or two-stage temperature range. Is.
また、特許文献3には、Cを0.2〜0.7質量%含有する鋼を、仕上温度(Ar3変態点-20℃)以上で熱間圧延した後、冷却速度120℃/秒超かつ冷却停止温度650℃以下で冷却を行い、次いで巻取温度600℃以下で巻取、酸洗後、焼鈍温度600℃以上Ac1変態点以下で焼鈍し、冷圧率30%以上で冷間圧延を行い、焼鈍温度600℃以上Ac1変態点以下で焼鈍するというものである。金属組織については、炭化物平均粒径を0.1μm以上2.0μm未満、炭化物を含まないフェライト粒の体積率を15%以下に制御することを特徴としている。
しかしながら、特許文献１及び特許文献２に記載の技術では、フェライト組織が初析フェライトからなり、初析フェライトと球状化炭化物を含むフェライトでは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。また、伸びフランジ性に関しても、打抜き加工時に打抜き端面近傍では、初析フェライトと球状化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、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するため、伸びフランジ性が劣化する。 However, in the techniques described in Patent Document 1 and Patent Document 2, the ferrite structure is composed of pro-eutectoid ferrite, and the amount of deformation is greatly different between pro-eutectoid ferrite and ferrite containing spheroidized carbides. For this reason, stress concentrates 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, resulting in deterioration of ductility. In addition, regarding stretch flangeability, in the vicinity of the punched end face during punching, the amount of deformation differs greatly between ferrites containing proeutectoid ferrite and spheroidized carbide, and voids are generated at the interface between the spheroidized structure and proeutectoid ferrite. Since this grows into a crack, stretch flangeability deteriorates.
この対策として、球状化焼鈍を強化することにより、全体として軟質化させることが考えられる。しかし、その場合は製造工程が長くなり、コストが増大するという問題がある。 As a countermeasure against this, it is conceivable to soften the whole by strengthening the spheroidizing annealing. However, in this case, there is a problem that the manufacturing process becomes long and the cost increases.
特許文献３に記載の技術では、冷却後に変態発熱を生じて温度が上昇し、初析フェライトの析出およびパーライト変態が進行し、炭化物の粗大化や不均一分散を生じ、特性の劣化を招きやすい。 In the technology described in Patent Document 3, transformation heat generation occurs after cooling, the temperature rises, precipitation of pro-eutectoid ferrite and pearlite transformation progress, and coarsening and non-uniform dispersion of carbides are likely to occur, resulting in deterioration of characteristics. .
また、最近では従来にもまして、生産性の向上の観点から、加工レベルに対する要求が厳しくなっている。そのため、高炭素鋼についても、加工度の増加等により、延性および穴拡げ性に優れた材料が求められている。 In recent years, demands for processing levels have become stricter from the viewpoint of improving productivity. Therefore, high carbon steel is also required to have excellent ductility and hole expansibility due to an increase in workability.
本発明は、かかる事情に鑑み、長時間を要する多段階焼鈍を用いることなく製造できる、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。 In view of such circumstances, an object of the present invention is to provide a high carbon cold-rolled steel sheet excellent in workability that can be manufactured without using multi-stage annealing that requires a long time.
本発明は、高炭素鋼板の延性および伸びフランジ性に及ぼす組成およびミクロ組織の影響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、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は、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大きな影響を及ぼしていることを見出した。 The present invention was made in the course of diligent research on the effects of composition and microstructure on the ductility and stretch flangeability of high carbon steel sheets. And in the process, it discovered that the factor which has a big influence on the ductility and stretch flangeability of a steel plate had a great influence not only on a composition and the shape and quantity of a carbide | carbonized_material but the dispersion state of the carbide | carbonized_material.
さらに、炭化物の形状としては炭化物平均粒径、炭化物の分散状態として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、それぞれ制御することにより、高炭素鋼板の延性および伸びフランジ性が向上することがわかった。 Furthermore, by controlling the average particle size of the carbide as the shape of the carbide, and the volume fraction of ferrite grains in which the volume fraction of the intragranular carbide is 10% or less as the carbide dispersion state, the ductility and elongation of the high carbon steel sheet are controlled. It was found that the flangeability was improved.
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、加工性に優れた高炭素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 steel plate excellent in workability 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：0.10〜0.35%、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的不純物からなり、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。
［２］前記［１］において、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。
［３］前記［１］または［２］のいずれかに記載の組成を有する鋼を、(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、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、600℃以上Ac1変態点以下の温度で再結晶焼鈍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。
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: 0.10 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, Cr: 0.05 to 0.30% is contained, the balance is made of iron and inevitable impurities, the ferrite average particle size is 2.0 μm or more, the carbide average particle size is 0.10 μm or more and less than 2.0 μm, and the volume fraction of intragranular carbide is A high carbon cold-rolled steel sheet excellent in workability, characterized by having a structure in which a volume fraction of ferrite grains of 10% or less is 50% or more.
[2] In the above [1], one or more 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% are contained by mass%. A high carbon cold-rolled steel sheet with excellent workability characterized by
[3] Hot-rolling the steel having the composition described in either [1] or [2] above at a finishing temperature of (Ar3 transformation point -20 ° C) or higher, and then cooling at a temperature exceeding 120 ° C / second 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 more and below the Ac1 transformation point, cold rolling at a cold pressure rate of 30% or more and recrystallization annealing at a temperature of 600 ° C or more and below the Ac1 transformation point A method for producing a high carbon cold-rolled steel sheet having excellent properties.
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 the present specification, “%” indicating the component of steel is “% by mass”.
本発明では、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。 In the present invention, not only the control of the component composition and manufacturing conditions, but also the ferrite particle size, carbide particle size, and the dispersion state of the carbide are controlled, thereby suppressing the generation of voids in press molding and hole expansion processing, and cracking. Can slow down the growth. As a result, it is possible to provide a high carbon cold-rolled steel sheet having excellent workability.
本発明の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、複雑形状の成形を少ない工程で成形でき、製造工程の省略による低コストでの製造が可能となる。 By using the high-carbon cold-rolled steel sheet of the present invention, it is possible to increase the degree of processing in processing of transmission parts and the like typified by gears. The manufacturing at a low cost is possible by omitting.
本発明の高炭素冷延鋼板は、下記に示す成分組成に制御し、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し、これらは本発明において最も重要な要件である。このように成分組成と金属組織（フェライト平均粒径）、炭化物の形状（炭化物平均粒径、）および炭化物の分散状態（粒内炭化物の体積率が10%以下であるフェライト粒）を規定し、全てを満足することにより、加工性に優れた高炭素冷延鋼板を得ることができる。 The high carbon cold-rolled steel sheet of the present invention is controlled to have the following component composition, the ferrite average particle size is 2.0 μm or more, the carbide average particle size is 0.10 μm or more and less than 2.0 μm, and the volume fraction of intragranular carbide is 10% or less. It is characterized by having a structure in which the volume fraction of ferrite grains is 50% or more, and these are the most important requirements in the present invention. Thus, the component composition and metal structure (ferrite average particle size), the shape of carbide (carbide average particle size), and the dispersion state of carbide (ferrite particles in which the volume fraction of intragranular carbide is 10% or less) By satisfying all, a high carbon cold-rolled steel sheet having excellent workability can be obtained.
そして、上記高炭素冷延鋼板は、(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、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、600℃以上Ac1変態点以下の温度で再結晶焼鈍することにより製造することが可能となる。このように、熱間圧延後、1次冷却、2次冷却、巻取および焼鈍までの条件をトータルで制御することにより、本発明の目的が達成される。 The high carbon cold-rolled steel sheet is hot-rolled at a finishing temperature of (Ar3 transformation point -20 ° C) or higher, and then at a cooling rate exceeding 120 ° C / second to a cooling stop temperature of 500 ° C or higher and 650 ° C or lower. After primary cooling, and 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 spheroidizing annealing at a temperature of 600 ° C or higher and lower than the Ac1 transformation point Then, it can be manufactured by performing cold rolling at a cold pressure ratio of 30% or more and performing recrystallization annealing at a temperature not lower than 600 ° C. and not higher than the Ac1 transformation point. Thus, the object of the present invention is achieved by controlling the conditions from hot rolling to primary cooling, secondary cooling, winding and annealing in total.
以下、本発明を詳細に説明する。 Hereinafter, the present invention will be described in detail.
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%未満の鋼では、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。
(2)Si：0.10〜0.35%
Siは、焼入れ性を向上させる元素である。Siが0.10%未満では焼入れ時の硬さが不足し、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0.10％以上0.35%以下、好ましくは0.10%以上0.30%以下とする。
(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。
(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、0.03%以下、好ましくは0.02%以下とする。
(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。
(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、0.08%以下とす
る。
(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、0.01%以下とする。
(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。
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. Steel with a C content of less than 0.2% cannot provide sufficient quenching hardness when applied 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: 0.10 to 0.35%
Si is an element that improves hardenability. If Si is less than 0.10%, the hardness at the time of quenching will be insufficient.On the other hand, if Si exceeds 0.35%, ferrite will harden due to solid solution strengthening, and ductility and stretch flangeability will deteriorate, causing cracks during molding processing. It becomes. Therefore, from the viewpoint of providing a steel sheet having both appropriate quenching hardness and workability, the Si content is set to 0.10% to 0.35%, preferably 0.10% to 0.30%.
(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% or less
P segregates at grain boundaries and deteriorates ductility and toughness, so 0.03% or less, preferably 0.02% or less.
(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% or less
If Al is added in excess, a large amount of AlN precipitates and lowers the hardenability, so 0.08% or less.
(7) N: 0.01% or less
If N is excessively contained, the ductility is lowered, so the content is made 0.01% or less.
(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.
また、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。 In addition, the steel of the present invention can achieve the desired characteristics with the above-mentioned essential additive elements, but in addition to the above-mentioned essential additive elements, it suppresses the formation of proeutectoid ferrite during hot rolling cooling, and improves the hardenability. , Mo, Ti, Nb may be added singly 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。不可避的不純物として、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は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。
(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未満では、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。
(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。さらに、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%以上とすることで、鋼板硬さが低下し、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は50%以上とする。
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 average ferrite particle diameter (average particle diameter of ferrite grains) is smaller 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) 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. When the carbide becomes finer, the generation of voids can be suppressed. However, when the average particle size of the carbide is less than 0.10 μm, the ductility decreases with increasing hardness 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) The volume fraction of ferrite grains with a volume fraction of intra-grain carbide of 10% or less is 50% or more. The dispersion state of carbide greatly affects the workability in general and the generation of voids in hole expansion processing. 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. Furthermore, 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, void formation and void connection are suppressed, ductility and stretch flange The characteristics are 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%以下であれば、ボイドの発生および連結を抑制する効果は十分に得られ、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場合と実質的に同じとみなすことができる。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, if the volume fraction of intragranular carbide is 10% or less, the effect of suppressing the generation and connection of voids can be sufficiently obtained, and there is no increase in hardness due to dispersion dispersion of carbide, ductility and stretch flangeability Can be considered to be substantially the same as when no intragranular carbide is contained. The dispersion state of carbides, that is, the volume fraction of ferrite grains in which the volume fraction of intragranular carbide is 10% or less is the production condition, particularly the primary cooling stop temperature after rolling, the secondary cooling holding temperature, the coiling temperature, And can be controlled by the 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 is determined by polishing the plate thickness section of the steel sheet sample, corroding it with a night, and observing about 3000 ferrite grains at 2000 times with a scanning electron microscope. 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、600℃以上Ac1変態点以下の温度で球状化焼鈍し、30%以上の冷圧率で冷間圧延を行い、600℃以上Ac1変態点以下の温度で再結晶焼鈍する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。
(1)仕上げ温度：(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化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。以上の理由により、(Ar3変態点-20℃)以上の仕上温度で仕上圧延することとする。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（1）で算出することができる。
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, followed by 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 600 ° C It is obtained by spheroidizing annealing at a temperature not lower than the Ac1 transformation point, cold rolling at a cold pressure ratio of 30% or higher, and recrystallization annealing at a temperature not lower than 600 ° C and not higher than the Ac1 transformation point. This will be described in detail below.
(1) Finishing temperature: (Ar3 transformation point -20 ° C) or higher If the finishing temperature when hot rolling the steel is less than (Ar3 transformation point -20 ° C), the ferrite transformation proceeds in part, so the pro-eutectoid ferrite Grains increase, voids are likely to occur at the interface between pro-eutectoid ferrite and ferrite containing spheroidized carbide, and ductility and stretch flangeability deteriorate. For the above reasons, finish rolling is performed at a finishing temperature of (Ar3 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 Ar3 transformation point (° C.) can be calculated by the following formula (1).
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。
Ar3 = 930.21-394.75C + 54.99Si-14.40Mn + 5.77Cr (1)
Here, the element symbol in a formula represents content (mass%) of each element.
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。
(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化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ー間隔が増大し、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。
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) Primary cooling rate: over 120 ° C / sec. When 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 prominently generated, voids are easily generated at the interface between pro-eutectoid ferrite and ferrite containing spheroidized carbides, and ductility and stretch flangeability deteriorate. Further, the interval between pearlite colonies and lamellar increases, leading to a long spheroidizing annealing time, resulting in an increase in 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.
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やパーライト等の析出物をより微細化し、加工性をより一層向上する上で好ましい。
(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、球状化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、次に行われる再結晶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、このような細粒効果および炭化物の微細分散効果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。
(4)2次冷却保持温度：500℃以上650℃以下
高炭素鋼板の場合、1次冷却停止温度後に、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、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、球状化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。これらの理由から、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。
(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。
(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。
(7)球状化焼鈍の焼鈍温度：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焼鈍を行う。球状化焼鈍の温度が600℃未満の場合、炭化物の球状化が不十分となり、焼鈍の効果が得られない。一方、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。以上の理由により、球状化焼鈍の温度は600℃以上Ac1変態点以下とする。なお、優れた加工性を得るには、焼鈍温度を680℃以上とすることが好ましい。また、Ac1変態点（℃）は次の式（2）で算出することができる。
Here, the cooling rate is an average cooling rate from the start of cooling after finish rolling to the stop of cooling. In addition, after finishing rolling, starting cooling within a period of more than 0.1 seconds and less than 1.0 seconds is necessary to further refine the precipitates such as ferrite crystal grains and pearlite after transformation and further improve workability. preferable.
(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 interval increases, and non-spherical carbides tend to remain after spheroidizing annealing. This non-spherical carbide is crushed during the cold rolling, and suppresses the recrystallization of ferrite during the subsequent recrystallization annealing to become fine grains. Further, most of the non-spherical carbides tend to be intragranular carbides. And with the strength increase by such a fine grain effect and the fine dispersion effect of a carbide, a ductility fall is caused and workability deteriorates. 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) Secondary cooling holding temperature: 500 ° C or higher and 650 ° C or lower In the case of high-carbon steel plates, the steel plate temperature may increase with the primary eutectoid ferrite transformation, pearlite transformation, and bainite transformation after the primary cooling stop temperature. Even if the primary cooling stop temperature is 650 ° C. or lower, if the temperature rises from the end of the primary cooling to the winding, ferrite is generated and the pearlite lamella spacing becomes coarse. Therefore, workability deteriorates due to the non-spherical carbide after spheroidizing annealing. Further, 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) 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) Pickling: Implementation 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) Annealing temperature of spheroidizing annealing: 600 ° C. or more and Ac1 transformation point or less After hot pickling the steel sheet, cold rolling is performed, but before that, annealing is performed to spheroidize the carbide. When the temperature of spheroidizing annealing is less than 600 ° C., the spheroidizing of the carbide becomes insufficient and the effect of annealing cannot be obtained. On the other hand, when the annealing temperature exceeds the Ac1 transformation point, a part is austenitized and pearlite is generated again during cooling, so that a spheroidized structure cannot be obtained. For the above reasons, the temperature of spheroidizing annealing is set to 600 ° C. or more and Ac1 transformation point or less. In order to obtain excellent workability, the annealing temperature is preferably 680 ° C. or higher. The Ac1 transformation point (° C.) can be calculated by the following formula (2).
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。
(8)冷間圧延の圧下率：30%以上
冷間圧延を行うことにより、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、再結晶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。よって、冷間圧延の圧下率は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。
(9)再結晶焼鈍の焼鈍温度：600℃以上Ac1変態点以下
冷間圧延後、再結晶および炭化物の球状化促進のために焼鈍を行う。焼鈍温度が600℃未満の場合、再結晶が不十分となり、未再結晶部が残存するため、加工性が劣化する。また、炭化物の球状化が不十分あるいは炭化物平均粒径が0.1μm未満となり、加工性が劣化する。一方、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、やはり加工性が劣化する。また、焼鈍温度を680℃以上とすることにより、炭化物の球状化率が高くなるため、高い伸び特性が得られる。また、より一層優れた加工性が得られる。以上の点から、より優れた加工性を得る場合には、焼鈍温度を680℃以上とすることが好ましい。また、Ac1変態点（℃）は上記式（2）で算出することができる。
Ac1 = 754.83-32.25C + 23.32Si-17.76Mn + 17.13Cr (2)
Here, the element symbol in a formula represents content (mass%) of each element.
(8) Cold rolling reduction: 30% or more By performing cold rolling, ferrite recrystallization during 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 unrecrystallized parts remain after recrystallization annealing, which deteriorates workability. 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) Annealing temperature of recrystallization annealing: 600 ° C. or more and Ac1 transformation point or less After cold rolling, annealing is performed to promote recrystallization and spheroidization of carbides. When the annealing temperature is less than 600 ° C., recrystallization becomes insufficient and unrecrystallized parts remain, so that workability deteriorates. Further, the spheroidization of the carbide is insufficient or the average particle size of the carbide is less than 0.1 μm, and the workability is deteriorated. On the other hand, when the annealing temperature exceeds the Ac1 transformation point, part of it becomes austenite and pearlite is generated again during cooling, so that the workability deteriorates. Further, by setting the annealing temperature to 680 ° C. or higher, the spheroidization rate of the carbide is increased, so that high elongation characteristics can be obtained. Further, further excellent processability can be obtained. From the above points, in order to obtain more excellent workability, the annealing temperature is preferably set to 680 ° C. or higher. The Ac1 transformation point (° C.) can be calculated by the above formula (2).
なお、本発明の高炭素鋼の成分調整には、転炉あるいは電気炉のどちらでも使用可能である。このように成分調整された高炭素鋼を、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。 It should be noted that either a converter or an electric furnace can be used to adjust the components of the high carbon steel of the present invention. The high carbon steel whose components have been adjusted in this way is made into a steel slab that 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.
再結晶焼鈍の後、必要に応じて調質圧延を行う。この調質圧延については焼入れ性には影響を及ぼさないことから、その条件に対して特に制限はない。 After recrystallization annealing, temper rolling is performed as necessary. Since this temper rolling does not affect the hardenability, there is no particular limitation on the conditions.
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部分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。 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 influenced 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.
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。 Thus, not only control of manufacturing conditions but also control of the average particle size of carbides and the dispersion state of carbides can suppress void connection and growth.
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、冷間圧延、および焼鈍を行い、板厚3.0mmの冷延鋼板を製造した。ここで、鋼板No.1〜6は製造条件が本発明範囲内の本発明例であり、鋼板No.7〜15は製造条件が本発明範囲から外れる比較例、鋼板No.16、17は鋼成分が本発明範囲から外れる比較例である。 A continuous cast slab of steel having the chemical composition shown in Table 1 is heated to 1250 ° C and hot rolled, cold rolled, and annealed under the conditions shown in Table 2 to produce a cold rolled steel sheet with a thickness of 3.0 mm did. Here, steel plates Nos. 1 to 6 are examples of the present invention in which the manufacturing conditions are within the scope of the present invention, steel plates No. 7 to 15 are comparative examples in which the manufacturing conditions are outside the scope of the present invention, and steel plates No. 16 and 17 are steel. This is a comparative example in which the components are out of the scope of 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。
＜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。
＜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。
＜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。
＜硬度＞
試料の切断面をバフ研磨仕上げ後、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。
＜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。
Next, a sample is taken from the cold-rolled steel sheet obtained as described above, and the average ferrite particle size, average carbide particle size and dispersion state are measured, and for performance evaluation, hardness, elongation (JIS5C direction), and elongation Flangeability was measured. Each measuring method and conditions are as follows.
<Ferrite average particle size>
From the light microscopic structure in the plate thickness section of the sample, it was cut by the cutting method described in JIS G 0552.
<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.
<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. .
<Hardness>
The cut surface of the sample was buffed and polished, and the Vickers hardness (Hv) was measured at the center of the plate thickness under a load of 500 gf.
<Elongation (JIS5 C direction)>
After cutting a JIS5 test piece in the direction of 90 ° to the rolling direction, perform a tensile test under the conditions of a crosshead of 10 mm / min and a distance between gauge points of 10 = 50 mm. Measurement was made to obtain an elongation amount El defined by the following equation.
El=100×(l-l0)/l0
＜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。
El = 100 × (ll 0 ) / l 0
<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×(d1-d0)/d0
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。
λ = 100 × (d 1 -d 0 ) / d 0
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.1〜6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%以上の発明例である。 In Table 3, the production conditions of the steel plates No. 1 to 6 are within the scope of the present invention, the average particle size of ferrite is 2.0 μm or more, the average particle 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 invention example in which the volume fraction of ferrite grains having a percentage of 50% or less is 50% or more.
一方、鋼板No.7〜15は製造条件が本発明範囲を外れた比較例、鋼板No.16、17は鋼成分が本発明から外れる比較例であり、鋼板No.9、11、14はフェライト平均粒径が2.0μm未満であり、鋼板No.7は炭化物平均粒径が下限0.10μm未満、鋼板No.8〜10、12〜17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が50%未満であり、いずれも本発明の範囲外である。 On the other hand, steel plates No. 7 to 15 are comparative examples in which the manufacturing conditions are outside the scope of the present invention, steel plates No. 16 and 17 are comparative examples in which the steel components are out of the present invention, and steel plates No. 9, 11, and 14 are ferrites. The average grain size is less than 2.0 μm, the steel plate No. 7 has a carbide average grain size of less than 0.10 μm lower limit, the steel plates No. 8 to 10 and 12 to 17 are ferrite grains having a volume fraction of intragranular carbide of 10% or less Is less than 50%, both of which are outside the scope of the present invention.
表3より、本発明例1〜6は、比較例7〜15に比べて、それぞれ同じ鋼種において、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。 From Table 3, Invention Examples 1-6 are improved in both ductility (El) and hole expansion rate (λ) in the same steel type as compared with Comparative Examples 7-15, and have excellent workability. I understand.
本発明の冷延鋼板は、自動車部品以外にも、延性(El)、優れた加工性（伸び及び伸びフランジ性）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 suitable not only for automobile parts but also for applications that require ductility (El) and excellent workability (elongation and stretch flangeability).
Claims (3)
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。 In mass%, C: 0.2 to 0.7%, Si: 0.10 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 : 0.05 to 0.30%, the balance consists of iron and inevitable impurities,
The ferrite average particle size is 2.0 μm or more, the carbide average particle size is 0.10 μm or more and less than 2.0 μm, the volume fraction of the intragranular carbide is 10% or less, and the volume fraction of ferrite grains is 50% or more. High carbon cold-rolled steel sheet with excellent workability.
(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、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、
600℃以上Ac1変態点以下の温度で再結晶焼鈍する
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。
Steel having the composition according to claim 1 or 2,
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,
Winding at a temperature of 600 ℃ or less,
After pickling, after spheroidizing annealing at a temperature of 600 ° C or higher and below Ac1 transformation point, cold rolling at a cold pressure rate of 30% or more,
A method for producing a high carbon cold-rolled steel sheet having excellent workability, characterized by recrystallization annealing at a temperature of 600 ° C. or more and below the Ac1 transformation point.
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