JP5011846B2 - High carbon hot rolled steel sheet and manufacturing method thereof - Google Patents
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Description
本発明は、加工性に優れた高炭素熱延鋼板およびその製造方法に関する。 The present invention relates to a high carbon hot-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、種々の複雑な形状に加工されるため優れた加工性がユーザーから求められる。一方、近年、部品製造コスト低減の要求が強くなり、加工工程の省略や加工方法の変更が行なわれている。例えば、非特許文献1に記載されているように、高炭素鋼板を用いた自動車駆動系部品の成形技術として、増肉成形を可能にし、大幅な工程短縮を実現した複動成形技術が開発され、一部実用化されている。それとともに、高炭素鋼板には、加工性に対する要求が益々強くなっており、より高い延性が求められている。また、部品によっては、打抜き加工後に穴拡げ加工(バーリング)を受ける場合が多いので、伸びフランジ性に優れていることも望まれている。さらに、歩留り向上にともなうコスト低減の観点から、鋼板の材質均一性も強く要望されている。特に、鋼板の板厚方向で表層部と中心部の硬度差が大きいと打抜き加工における打抜き工具の劣化が激しくなるので、板厚方向の硬度均一性が切望されている。 High carbon steel sheets used for tools or automobile parts (gears, missions) and the like are processed into various complicated shapes, and thus excellent workability is required from users. On the other hand, in recent years, demands for reducing component manufacturing costs have increased, and processing steps have been omitted and processing methods have been changed. For example, as described in Non-Patent Document 1, a double-acting molding technology has been developed as a molding technology for automobile drive system parts using high-carbon steel sheets that enables thickening molding and realizes a significant process shortening. Some have been put to practical use. At the same time, high carbon steel sheets are increasingly demanded for workability and are required to have higher ductility. In addition, some parts are often subjected to hole expansion processing (burring) after punching, and therefore, it is also desired that they have excellent stretch flangeability. Furthermore, from the viewpoint of cost reduction accompanying yield improvement, there is a strong demand for material uniformity of the steel sheet. In particular, if the hardness difference between the surface layer portion and the center portion in the plate thickness direction of the steel plate is large, the punching tool in the punching process is severely deteriorated. Therefore, hardness uniformity in the plate thickness direction is desired.
こうした要求に答えるべく、高炭素鋼板の加工性や材質均一性を向上させるために、従来からいくつかの技術が検討されている。例えば、特許文献1には、ホットランテーブルを加速冷却ゾーンと空冷ゾーンに2分割し、仕上圧延後の鋼帯を冷却ゾーンの長さ、鋼板の搬送速度、化学成分などで決まる特定の温度以下に加速冷却した後空冷して、コイル長手方向の材質均一性に優れる高炭素鋼帯を安定して製造する方法が提案されている。なお、同公報における加速冷却域での冷却速度は、第8図から20〜30℃/秒程度である。また、特許文献2には、所定の化学成分の高炭素鋼を熱間圧延し、脱スケールを行った後、95容量%以上の水素雰囲気中で、化学成分で規定された加熱速度や均熱時間の条件で焼鈍後、100℃/ｈｒ以下の冷却速度で冷却して、軟質で、組織の均一性や加工性に優れた高炭素鋼帯を製造する方法が提案されている。さらに、特許文献3には、(Ac1変態点+30℃)以上の仕上温度で圧延された鋼板を10〜100℃/秒の冷却速度で20〜500℃の温度まで冷却し、1〜10秒保持後、500〜(Ac1変態点+30℃)の温度域に再加熱して巻取り、必要に応じて650〜(Ac1変態点+30℃)で1時間以上均熱することにより加工性の良好な高炭素薄鋼板を製造する方法が提案されている。さらにまた、特許文献4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℃以下で巻取り、焼鈍温度640℃以上Ac1変態点以下で焼鈍することにより、伸びフランジ性に優れた高炭素熱延鋼板を製造する方法が提案されている。
しかしながら、これらの従来技術はいずれも、板厚方向まで含めた材質の均一性を確保するものではなく、またこのような均一性と伸びフランジ性を両立させるものではなかった。さらに、これらの従来技術には以下のような問題もある。 However, none of these prior arts ensure the uniformity of the material including the thickness direction, and do not achieve such uniformity and stretch flangeability. Furthermore, these conventional techniques also have the following problems.
特許文献1に記載の方法では、製造されるのが熱間圧延ままの鋼板であるため、必ずしも優れた伸びや伸びフランジ性が得られるとは限らない。 In the method described in Patent Document 1, since it is a hot-rolled steel sheet that is produced, excellent elongation and stretch flangeability are not always obtained.
特許文献2に記載の方法では、熱延条件によっては初析フェライトとラメラー状の炭化物を有するパーライトからなるミクロ組織が形成され、その後の焼鈍でラメラー状の炭化物が微細な球状化炭化物となる。この微細な球状化炭化物は穴拡げ加工時にボイド発生の起点になり、発生したボイドが連結して破断を誘発するため、優れた伸びフランジ性が得られない。
In the method described in
特許文献3に記載の方法では、熱間圧延後の鋼板を所定の条件で冷却後、直接通電法などで再加熱しているため特別な設備が必要となるばかりか、膨大な電力エネルギーが必要となる。また、再加熱後に巻取った鋼板には微細な球状化炭化物が形成されるため、上記と同様の理由で優れた伸びフランジ性が得られない場合が多い。 The method described in Patent Document 3 requires not only special equipment but also enormous power energy because the steel sheet after hot rolling is cooled under a predetermined condition and then reheated by a direct current method or the like. It becomes. In addition, since fine spheroidized carbide is formed on the steel sheet wound after reheating, excellent stretch flangeability is often not obtained for the same reason as described above.
本発明は、伸びフランジ性と板厚方向の硬度均一性に優れた高炭素熱延鋼板およびその製造方法を提供することを目的とする。 An object of the present invention is to provide a high carbon hot-rolled steel sheet excellent in stretch flangeability and hardness uniformity in the sheet thickness direction, and a method for producing the same.
本発明者らは、高炭素熱延鋼板の伸びフランジ性および硬度に及ぼすミクロ組織の影響について鋭意研究を進めた結果、製造条件、特に、熱間圧延後の冷却条件、巻取温度、および焼鈍温度を適切に制御することが極めて重要であることを見出した。そして、後述する測定法で求めた粒径が0.5μm未満の炭化物の全炭化物に対する体積率を15%以下に制御することにより、伸びフランジ性が向上し、板厚方向の硬度が均一になることを見出した。また、さらに厳密に熱間圧延後の冷却条件、巻取温度を制御し、炭化物の前記体積率を10%以下に制御することにより、より優れた伸びフランジ性および硬度分布の均一性が得られることを見出した。 As a result of diligent research on the influence of microstructure on stretch flangeability and hardness of high-carbon hot-rolled steel sheets, the present inventors have determined that manufacturing conditions, particularly cooling conditions after hot rolling, coiling temperature, and annealing, We have found that it is very important to control the temperature appropriately. And, by controlling the volume ratio of carbides with a particle size of less than 0.5 μm, determined by the measurement method described below, to 15% or less, the stretch flangeability is improved and the hardness in the thickness direction becomes uniform. I found. In addition, by controlling the cooling conditions and coiling temperature after hot rolling more strictly and controlling the volume fraction of carbide to 10% or less, more excellent stretch flangeability and hardness distribution uniformity can be obtained. I found out.
本発明は、以上の知見に基づいてなされたものであり、C:0.2〜0.7質量%、Si:2質量%以下、Mn:2質量%以下、P:0.03質量%以下、S:0.03質量%以下、Sol.Al:0.08質量%以下、N:0.01質量%以下を含有し、残部Feおよび不可避的不純物からなる組成の鋼を、(Ar3変態点-20℃)以上の仕上温度で熱間圧延して熱延鋼板とする工程と、前記熱延鋼板を、80℃/秒以上120℃/秒未満の冷却速度で540℃以上600℃以下の温度(以後、冷却停止温度と呼ぶ)まで冷却する工程と、前記冷却後の熱延鋼板を、490℃以上550℃以下の巻取温度で巻取る工程と、前記巻取り後の熱延鋼板を、炭化物の球状化のために、640℃以上Ac1変態点以下の焼鈍温度で8時間以上80時間以下焼鈍(以後、熱延鋼板焼鈍と呼ぶ)する工程とを有する高炭素熱延鋼板の製造方法を提供する。 The present invention has been made based on the above findings, C: 0.2-0.7% by mass, Si: 2% by mass or less, Mn: 2% by mass or less, P: 0.03% by mass or less, S: 0.03% by mass Below, Sol.Al: 0.08% by mass or less, N: 0.01% by mass or less , steel of the composition consisting of the balance Fe and unavoidable impurities , hot at a finishing temperature of (Ar 3 transformation point -20 ° C.) or more The process of rolling into a hot-rolled steel sheet and cooling the hot-rolled steel sheet to a temperature of 540 ° C. or higher and 600 ° C. or lower (hereinafter referred to as a cooling stop temperature) at a cooling rate of 80 ° C./second or higher and less than 120 ° C./second. A step of winding the hot-rolled steel sheet after cooling at a winding temperature of 490 ° C. or higher and 550 ° C. or lower , and a hot-rolled steel sheet after the winding for spheroidizing carbides at 640 ° C. or higher. There is provided a method for producing a high carbon hot-rolled steel sheet comprising a step of annealing at an annealing temperature not higher than the Ac 1 transformation point for not less than 8 hours and not more than 80 hours (hereinafter referred to as hot-rolled steel sheet annealing).
本発明は、また、熱延球状化焼鈍材である高炭素熱延鋼板であって、C:0.2〜0.7質量%、Si:2質量%以下、Mn:2質量%以下、P:0.03質量%以下、S:0.03質量%以下、Sol.Al:0.08質量%以下、N:0.01質量%以下を含有し、残部Feおよび不可避的不純物からなる組成を有し、粒径0.5μm未満の炭化物の体積率が全炭化物に対する体積率で10%以下であり、かつ板厚方向における最大硬度Hvmaxと最小硬度Hvminの差ΔHv(=Hvmax-Hvmin)が8以下である、高炭素熱延鋼板を提供する。 The present invention is also a high-carbon hot-rolled steel sheet that is a hot-rolled spheroidized annealing material, C: 0.2 to 0.7% by mass, Si: 2% by mass or less, Mn: 2% by mass or less, P: 0.03% by mass Hereinafter, S: 0.03% by mass or less, Sol.Al: 0.08% by mass or less, N: 0.01% by mass or less, having a composition composed of the balance Fe and inevitable impurities, the volume of carbide having a particle size of less than 0.5 μm Provided is a high carbon hot-rolled steel sheet having a rate of 10% or less in terms of the volume ratio relative to the total carbides, and a difference ΔHv (= Hvmax−Hvmin) between the maximum hardness Hvmax and the minimum hardness Hvmin in the sheet thickness direction being 8 or less.
さらに、鋼の組成には、上記組成に加えて、次の含有量の範囲のCr、Moのうちから選ばれた少なくとも1種を含有させることも可能である；
Cr:3.5質量%以下、Mo:0.7質量%以下。
Furthermore, in addition to the above composition, the steel composition may contain at least one selected from Cr and Mo in the following content range;
Cr: 3.5% by mass or less, Mo: 0.7% by mass or less .
本発明により、特別な設備を必要とせずに、伸びフランジ性と板厚方向の硬度均一性がともに優れた高炭素熱延鋼板を製造できるようになった。 According to the present invention, it is possible to produce a high carbon hot-rolled steel sheet excellent in both stretch flangeability and hardness uniformity in the sheet thickness direction without requiring special equipment.
以下に、本発明である高炭素熱延鋼板およびその製造方法について詳細に説明する。
＜鋼組成＞
1)C量
Cは、炭化物を形成し、焼入後の硬度を付与する重要な元素である。C量が0.2質量%未満では、熱間圧延後に初析フェライトの生成が顕著となり、熱延板焼鈍後の粒径が0.5μm未満の炭化物の体積率が増加し、伸びフランジ性や板厚方向の硬度均一性が劣化する。その上、焼入後も機械構造用部品としての十分な強度が得られない。一方、C量が0.7質量%を超えると、たとえ粒径が0.5μm未満の炭化物の体積率が15%以下であっても十分な伸びフランジ性が得られない。また、熱間圧延後の硬度が著しく高くなり、鋼板が脆くなるため取扱いに不便となるばかりか、焼入後の機械構造用部品としての強度も飽和する。したがって、C量は0.2〜0.7質量%に規定する。なお、焼入れ後の硬度をより重視する場合は、C量は0.5質量%超えに、また、加工性をより重視する場合は、C量は0.5質量%以下とすることが好ましい。
Below, the high carbon hot-rolled steel sheet and its manufacturing method which are this invention are demonstrated in detail.
<Steel composition>
1) C amount
C is an important element that forms a carbide and imparts hardness after quenching. When the amount of C is less than 0.2% by mass, the formation of proeutectoid ferrite becomes prominent after hot rolling, the volume fraction of carbides with a grain size of less than 0.5μm after hot-rolled sheet annealing increases, stretch flangeability and thickness direction. The hardness uniformity of is deteriorated. In addition, sufficient strength as a machine structural component cannot be obtained even after quenching. On the other hand, if the amount of C exceeds 0.7% by mass, sufficient stretch flangeability cannot be obtained even if the volume fraction of the carbide having a particle size of less than 0.5 μm is 15% or less. In addition, the hardness after hot rolling becomes extremely high and the steel sheet becomes brittle, which is inconvenient to handle, and the strength as a machine structural part after quenching is saturated. Therefore, the amount of C is defined as 0.2 to 0.7% by mass. When the hardness after quenching is more important, the C amount is preferably more than 0.5% by mass. When the workability is more important, the C amount is preferably 0.5% by mass or less.
C以外のその他の元素については、特に、規定しないが、Mn、Si、P、S、Sol.Al、Nなどの元素を通常の範囲で含有させることができる。しかし、Siは、炭化物を黒鉛化し、焼入性を阻害する傾向があるので2質量%以下に、Mnは、過剰の添加は延性の低下を引き起こす傾向があるので2質量%以下に、P、Sは、過剰に含有すると延性が低下し、またクラックも生成しやすくなるのでともに0.03質量%以下に、Sol.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ので0.08質量%以下に、Nは、過剰に含有すると延性が低下するので0.01質量%以下にすることが望ましい。好ましくは、それぞれSi:0.5質量%以下、Mn:1質量%以下、P:0.02質量%以下、S:0.01質量%以下、Sol.Al:0.05質量%以下、N:0.005質量%以下である。ここで、これらの各元素を所定量以下、例えば0.0001質量%未満に低減するにはコスト増を招くので、0.0001質量%以上程度の含有は許容することが好ましい。なお、P、S、Nの含有量は、上記の目的のため、極力低減することがより好ましい。また、特にSを低減することは伸びフランジ性の改善にも効果的であり、この観点からは、Sは0.007質量%以下まで低減することが好ましい。より好ましくは0.0045質量%以下である。なお、例えば、Mnは、固溶強化により鋼の強度を増加するとともに、焼入れ性向上の目的で、Siは、脱酸剤として作用するとともに、固溶強化により強度(硬さ)を増加させる目的で、Alは、脱酸剤として作用するとともに、Nと結合してAlNを形成し、オーステナイト粒の粗大化防止の目的で、Mnは0.2質量%以上、Siは0.01質量%以上、AlはSol.Alで0.015質量%以上含有することが好ましい。 Other elements other than C are not particularly specified, but elements such as Mn, Si, P, S, Sol. Al, N, and the like can be contained in a normal range. However, Si graphitizes carbides and tends to inhibit hardenability, so that it is 2% by mass or less, and Mn tends to cause deterioration in ductility, so that Mn is less than 2% by mass, P, If S is added excessively, the ductility decreases and cracks are likely to be generated, so both are 0.03% by mass or less, and if added too much, AlN precipitates a large amount and lowers the hardenability. If N is contained excessively, the ductility decreases when it is excessively contained, so it is desirable to make it 0.01% by mass or less. Preferably, Si: 0.5% by mass or less, Mn: 1% by mass or less, P: 0.02% by mass or less, S: 0.01% by mass or less, Sol.Al: 0.05% by mass or less, and N: 0.005% by mass or less. Here, in order to reduce each of these elements to a predetermined amount or less, for example, less than 0.0001% by mass, an increase in cost is caused. Therefore, it is preferable to allow the content of about 0.0001% by mass or more. In addition, it is more preferable to reduce the content of P, S, and N as much as possible for the above purpose. In particular, reducing S is effective in improving stretch flangeability, and from this viewpoint, S is preferably reduced to 0.007% by mass or less. More preferably, it is 0.0045 mass% or less. For example, Mn increases the strength of steel by solid solution strengthening, and for the purpose of improving hardenability, Si acts as a deoxidizer, and increases the strength (hardness) by solid solution strengthening In addition, Al acts as a deoxidizer and combines with N to form AlN. For the purpose of preventing coarsening of austenite grains, Mn is 0.2 mass% or more, Si is 0.01 mass% or more, and Al is Sol. It is preferable to contain 0.015 mass% or more of .Al.
さらに、例えば、焼入れ性の向上や焼戻し軟化抵抗の向上を目的として、通常添加される範囲でB、Cr、Ni、Mo、Cu、Ti、Nb、W、V、Zr等の少なくとも一つの元素を添加しても本発明の効果が損なわれることはない。具体的には、これらの元素は、B:0.005質量%以下、Cr:3.5質量%以下、Ni:3.5質量%以下、Mo:0.7質量%以下、Cu:0.1質量%以下、Ti:0.1質量%以下、Nb:0.1質量%以下、W、V、Zr:合計で0.1質量%以下含有させることができる。この目的のためには、Bは0.0005質量%以上、Crは0.05質量%以上、Niは0.05質量%以上、Moは0.05質量%以上、Cuは0.01質量%以上、Tiは0.01質量%以上、Nbは0.01質量%以上、W、V、Zrは合計で0.01質量%以上含有させることが好ましい。 Further, for example, for the purpose of improving hardenability and temper softening resistance, at least one element such as B, Cr, Ni, Mo, Cu, Ti, Nb, W, V, Zr, etc. is added in a range that is usually added. Even if it adds, the effect of this invention is not impaired. Specifically, these elements are B: 0.005 mass% or less, Cr: 3.5 mass% or less, Ni: 3.5 mass% or less, Mo: 0.7 mass% or less, Cu: 0.1 mass% or less, Ti: 0.1 mass% Hereinafter, Nb: 0.1% by mass or less, W, V, Zr: 0.1% by mass or less in total can be contained. For this purpose, B is 0.0005% by mass or more, Cr is 0.05% by mass or more, Ni is 0.05% by mass or more, Mo is 0.05% by mass or more, Cu is 0.01% by mass or more, Ti is 0.01% by mass or more, Nb Is preferably 0.01% by mass or more, and W, V, and Zr are preferably contained in a total of 0.01% by mass or more.
上記以外の残部はFeおよび不可避的不純物とすることが好ましいが、さらにまた、製造過程でSn、Pb等の元素が不純物として混入しても本発明の効果には影響を及ぼさない。
＜製造条件＞
2)熱間圧延の仕上温度
仕上温度が(Ar3変態点-20℃)未満では、フェライト変態が部分的に進行するため粒径が0.5μm未満の炭化物の体積率が増加し、伸びフランジ性と板厚方向の硬度均一性が劣化する。したがって、熱間圧延の仕上温度は(Ar3変態点-20℃)以上とする。なお、Ar3変態点は次の式(1)から計算できるが、実際に測定した温度を用いてもよい。
Ar3変態点=910-203×[C]1/2+44.7×[Si]-30×[Mn] ・・・(1)
ここで、[M]は元素Mの含有量(質量%)を表す。なお、含有元素に応じて、補正項を導入してもよく、例えば、CrやMo、Niを含有する場合には、-11×[Cr]、+31.5×[Mo]、-15.2×[Ni]といった補正項を式(1)の右辺に加えてよい。
The balance other than the above is preferably made of Fe and inevitable impurities. Furthermore, even if elements such as Sn and Pb are mixed as impurities in the production process, the effect of the present invention is not affected.
<Production conditions>
2) Finishing temperature of hot rolling If the finishing temperature is less than (Ar 3 transformation point -20 ° C), the ferrite transformation proceeds partially, so the volume fraction of carbides with grain size less than 0.5μm increases and stretch flangeability And the hardness uniformity in the plate thickness direction deteriorates. Therefore, the finishing temperature of hot rolling is (Ar 3 transformation point −20 ° C.) or higher. The Ar 3 transformation point can be calculated from the following equation (1), but the actually measured temperature may be used.
Ar 3 transformation point = 910-203 × [C] 1/2 + 44.7 × [Si] -30 × [Mn] (1)
Here, [M] represents the content (mass%) of the element M. A correction term may be introduced depending on the contained elements. For example, when Cr, Mo, or Ni is contained, -11 × [Cr], + 31.5 × [Mo], −15.2 × [Ni ] May be added to the right side of equation (1).
3)熱間圧延後の冷却条件
熱間圧延後の冷却速度が60℃/秒未満であると、オーステナイトの過冷度が小さくなり、熱間圧延後に初析フェライトの生成が顕著となる。その結果、熱延鋼板焼鈍後の粒径が0.5μm未満の炭化物の体積率が15%を超え、伸びフランジ性と板厚方向の硬度均一性が劣化する。一方、冷却速度が120℃/秒を超える場合は、板厚方向で表層部と中央部の温度差が大きくなり、中央部において初析フェライトの生成が顕著となる。その結果、上記と同様に、伸びフランジ性と板厚方向の硬度均一性が劣化する。この傾向は、熱延鋼板の板厚が4.0mm以上となると特に顕著となる。すなわち、特に板厚方向の硬度を均一とするためには、適正な冷却速度があり、冷却速度が過大でも過小でも所望の硬度均一性を得ることができない。従来技術においては、特に冷却速度の適正化がなされていないため、硬度均一性が確保できないのである。したがって、熱間圧延後の冷却速度は60℃/秒以上120℃/秒未満とする。さらに、粒径が0.5μm未満の炭化物の体積率を10%以下とする場合は、冷却速度を80℃/秒以上120℃/秒未満とする。なお、冷却速度は115℃/秒以下とすることが、より好ましい。
3) Cooling conditions after hot rolling When the cooling rate after hot rolling is less than 60 ° C / second, the degree of supercooling of austenite becomes small, and the formation of proeutectoid ferrite becomes noticeable after hot rolling. As a result, the volume fraction of the carbide having a particle size of less than 0.5 μm after annealing of the hot-rolled steel sheet exceeds 15%, and the stretch flangeability and the hardness uniformity in the sheet thickness direction deteriorate. On the other hand, when the cooling rate exceeds 120 ° C./second, the temperature difference between the surface layer portion and the central portion increases in the thickness direction, and proeutectoid ferrite is prominently generated in the central portion. As a result, similarly to the above, stretch flangeability and hardness uniformity in the plate thickness direction deteriorate. This tendency is particularly remarkable when the thickness of the hot-rolled steel sheet is 4.0 mm or more. That is, in particular, in order to make the hardness in the plate thickness direction uniform, there is an appropriate cooling rate, and the desired hardness uniformity cannot be obtained even if the cooling rate is too high or too low. In the prior art, since the cooling rate is not particularly optimized, hardness uniformity cannot be ensured. Therefore, the cooling rate after hot rolling is set to 60 ° C./second or more and less than 120 ° C./second. Furthermore, when the volume ratio of the carbide having a particle size of less than 0.5 μm is 10% or less, the cooling rate is 80 ° C./second or more and less than 120 ° C./second. The cooling rate is more preferably 115 ° C./second or less.
こうした冷却速度によって冷却する熱延鋼板の終点温度、すなわち冷却停止温度が650℃より高いと、熱延鋼板を巻取るまでの冷却中に初析フェライトが生成するとともに、ラメラー状の炭化物を有するパーライトが生成する。その結果、熱延鋼板焼鈍後の粒径が0.5μm未満の炭化物の体積率が15%を超え、伸びフランジ性と板厚方向の硬度均一性が劣化する。したがって、冷却停止温度は650℃以下とする。冷却停止温度は600℃以下とすることがさらに好ましい。なお、粒径が0.5μm未満の炭化物の体積率を10%以下とする場合は、前記したように冷却速度を80℃/秒以上120℃/秒未満(好ましくは115℃/秒以下)とするとともに冷却停止温度を600℃以下とする。温度の測定精度上の問題があるので、冷却停止温度は500℃以上とすることが好ましい。なお、冷却停止温度に到達した後は、特に規定する必要がなく、自然冷却してもよいし、冷却力を弱めて強制冷却を継続してもよい。鋼板の均一性などの観点からは復熱を抑制する程度に強制冷却することが好ましい。 When the end point temperature of the hot-rolled steel sheet cooled by such a cooling rate, that is, the cooling stop temperature is higher than 650 ° C., proeutectoid ferrite is generated during cooling until the hot-rolled steel sheet is wound, and pearlite having lamellar carbides. Produces. As a result, the volume fraction of the carbide having a particle size of less than 0.5 μm after annealing of the hot-rolled steel sheet exceeds 15%, and the stretch flangeability and the hardness uniformity in the sheet thickness direction deteriorate. Therefore, the cooling stop temperature is set to 650 ° C. or lower. The cooling stop temperature is more preferably 600 ° C. or lower. When the volume ratio of the carbide having a particle size of less than 0.5 μm is 10% or less, the cooling rate is 80 ° C./second or more and less than 120 ° C./second (preferably 115 ° C./second or less) as described above. At the same time, the cooling stop temperature is set to 600 ° C. or lower. Since there is a problem with temperature measurement accuracy, the cooling stop temperature is preferably 500 ° C. or higher. It should be noted that after reaching the cooling stop temperature, there is no particular need to define, and natural cooling may be performed, or forced cooling may be continued by weakening the cooling power. From the viewpoint of the uniformity of the steel sheet, it is preferable to perform forced cooling to such an extent that recuperation is suppressed.
4)巻取温度
冷却後の熱延鋼板は巻取られるが、そのとき、巻取温度が600℃を超えるとラメラー状の炭化物を有するパーライトが生成する。その結果、熱延鋼板焼鈍後の粒径が0.5μm未満の炭化物の体積率が15%を超え、伸びフランジ性と板厚方向の硬度均一性が劣化する。したがって、巻取温度は600℃以下とする。なお、前記急冷の効果を十分に得るため、巻取温度は前記冷却停止温度よりも低温とすることが好ましい。さらに、粒径が0.5μm未満の炭化物の体積率を10%以下とする場合は、前記したように冷却速度を80℃/秒以上120℃/秒未満(好ましくは115℃/秒以下)とし、冷却停止温度を600℃以下とするとともに、巻取温度を550℃以下とする。なお、熱延鋼板の形状が劣化するため、巻取温度は200℃以上とすることが好ましく、350℃以上とすることがより好ましい。
4) Winding temperature The hot-rolled steel sheet after cooling is wound, but at that time, when the winding temperature exceeds 600 ° C, pearlite having lamellar carbides is generated. As a result, the volume fraction of the carbide having a particle size of less than 0.5 μm after annealing of the hot-rolled steel sheet exceeds 15%, and the stretch flangeability and the hardness uniformity in the sheet thickness direction deteriorate. Therefore, the coiling temperature is 600 ° C. or less. In order to sufficiently obtain the effect of the rapid cooling, the winding temperature is preferably lower than the cooling stop temperature. Furthermore, when the volume fraction of the carbide having a particle size of less than 0.5 μm is 10% or less, the cooling rate is 80 ° C./second or more and less than 120 ° C./second (preferably 115 ° C./second or less) as described above. The cooling stop temperature is set to 600 ° C. or lower, and the winding temperature is set to 550 ° C. or lower. Note that, since the shape of the hot-rolled steel sheet deteriorates, the winding temperature is preferably 200 ° C. or higher, and more preferably 350 ° C. or higher.
5)スケール除去(酸洗など)
巻取り後の熱延鋼板は、通常、次の熱延鋼板焼鈍を行う前にスケール除去される。スケ−ル除去手段は、特に制約はないが、通常の方法で酸洗することが好ましい。
5) Scale removal (pickling etc.)
The rolled hot-rolled steel sheet is usually scaled before performing the next hot-rolled steel sheet annealing. The scale removing means is not particularly limited, but is preferably pickled by a normal method.
6)熱延鋼板焼鈍の温度
酸洗などによりスケール除去した後の熱延鋼板は、炭化物の球状化を図るために熱延球状化焼鈍として熱延鋼板焼鈍が施される。そのとき、熱延鋼板焼鈍の温度が640℃未満では炭化物の球状化が不十分であったり、粒径が0.5μm未満の炭化物の体積率が増加し、伸びフランジ性および板厚方向の硬度均一性が劣化する。一方、焼鈍温度がAc1変態点を超えるとオーステナイト化が部分的に進行し、冷却中に再度パーライトが生成するため、伸びフランジ性および板厚方向の硬度均一性が劣化する。したがって、熱延鋼板焼鈍の温度は640℃以上Ac1変態点以下とする。より優れた伸びフランジ性を得るために、熱延鋼板焼鈍の温度を680℃以上とすることが好ましい。なお、Ac1変態点は次の式(2)から計算できるが、実際に測定した温度を用いてもよい。
Ac1変態点=754.83-32.25×[C]+23.32×[Si]-17.76×[Mn] ・・・(2)
ここで、[M]は元素Mの含有量(質量%)を表す。なお、含有元素に応じて、補正項を導入してもよく、例えば、CrやMo、Vを含有する場合には、+17.3×[Cr]、+4.51×[Mo]、+15.62×[V]といった補正項を式(2)の右辺に加えてよい。
6) Temperature of hot-rolled steel sheet annealing The hot-rolled steel sheet after scale removal by pickling or the like is subjected to hot-rolled steel sheet annealing as hot-rolled spheroidizing annealing in order to spheroidize carbides. At that time, if the temperature of annealing of the hot-rolled steel sheet is less than 640 ° C., the spheroidization of the carbide is insufficient, or the volume fraction of the carbide having a particle diameter of less than 0.5 μm increases, stretch flangeability and uniform hardness in the sheet thickness direction. Deteriorates. On the other hand, when the annealing temperature exceeds the Ac 1 transformation point, austenitization partially proceeds and pearlite is generated again during cooling, so that stretch flangeability and hardness uniformity in the thickness direction are deteriorated. Therefore, the temperature of the hot-rolled steel sheet annealing is set to 640 ° C. or higher and Ac 1 transformation point or lower. In order to obtain more excellent stretch flangeability, it is preferable to set the temperature of hot-rolled steel sheet annealing to 680 ° C or higher. The Ac 1 transformation point can be calculated from the following equation (2), but the actually measured temperature may be used.
Ac 1 transformation point = 754.83-32.25 × [C] + 23.32 × [Si] -17.76 × [Mn] (2)
Here, [M] represents the content (mass%) of the element M. A correction term may be introduced depending on the contained elements. For example, when Cr, Mo, or V is contained, + 17.3 × [Cr], + 4.51 × [Mo], + 15.62 × [V ] May be added to the right side of Equation (2).
なお、焼鈍時間は8〜80時間程度が好ましい。得られた鋼板中の炭化物は球状化し、平均のアスペクト比で約5.0以下となる(板厚の約1/4の位置で測定した値)。 The annealing time is preferably about 8 to 80 hours. The carbides in the obtained steel plate are spheroidized and have an average aspect ratio of about 5.0 or less (value measured at about 1/4 of the plate thickness).
本発明の高炭素鋼を溶製するには、転炉、電気炉どちらも使用可能である。また、こうして溶製された高炭素鋼は、造塊−分塊圧延または連続鋳造によりスラブとされる。スラブは、通常、加熱(再加熱)された後、熱間圧延される。なお、連続鋳造で製造されたスラブの場合は、そのままあるい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適用してもよい。また、スラブを再加熱して熱間圧延する場合は、スケールによる表面状態の劣化を避けるためにスラブ加熱温度を1280℃以下とすることが好ましい。熱間圧延は、粗圧延を省略して仕上圧延だけで行うこともできる。なお、仕上温度を確保するため、熱間圧延中にシートバーヒータ等の加熱手段により被圧延材の加熱を行ってもよい。また、球状化促進あるいは硬度低減のため、巻取り後にコイルを徐冷カバー等の手段で保温してもよい。熱延鋼板の板厚は、本発明の製造条件が維持できる限りにおいて特に制限はないが、1.0〜10.0mmの熱延鋼板が操業上特に好適である。 To melt the high carbon steel of the present invention, both a converter and an electric furnace can be used. Further, the high carbon steel thus melted is made into a slab by ingot-bundling rolling or continuous casting. The slab is usually heated (reheated) and then hot rolled. In addition, in the case of the slab manufactured by continuous casting, you may apply direct feed rolling which rolls as it is or in order to suppress a temperature fall. Further, when the slab is reheated and hot-rolled, the slab heating temperature is preferably 1280 ° C. or lower in order to avoid deterioration of the surface state due to the scale. Hot rolling can be performed only by finish rolling, omitting rough rolling. In order to secure the finishing temperature, the material to be rolled may be heated by a heating means such as a sheet 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. The thickness of the hot-rolled steel sheet is not particularly limited as long as the production conditions of the present invention can be maintained, but a hot-rolled steel sheet having a thickness of 1.0 to 10.0 mm is particularly suitable for operation.
熱延鋼板焼鈍は、箱焼鈍、連続焼鈍いずれでも行える。熱延鋼板焼鈍後は、必要に応じて調質圧延を行う。この調質圧延は焼入れ性に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。 Hot-rolled steel sheet annealing can be performed by either box annealing or continuous annealing. After hot-rolled steel sheet annealing, temper rolling is performed as necessary. Since this temper rolling does not affect the hardenability, there is no particular limitation on the conditions.
上記本発明の方法で製造された熱延鋼板は、熱延球状化焼鈍を施された熱延鋼板であり、上記したように、平均のアスペクト比が約5.0以下と、球状化された炭化物を有する熱延鋼板である。また、粒径0.5μm未満の炭化物の体積率が全炭化物の体積率の15%以下、より好ましくは10%以下であり、伸びフランジ性に優れる。ここで、粒径0.5μm未満といった微細炭化物の体積率を上記のように低減することにより、伸びフランジ性が改善されるのは、このような微細炭化物は穴拡げ加工時にボイド発生の起点となり、発生したボイドが連結して破断を誘発するが、この炭化物量を低減することによりボイド発生の起点を減少させることができるためと考えられる。さらに、後述する図1に示すように、粒径0.5μm未満といった微細炭化物の体積率を上記のように低減した本発明の鋼板は、板厚方向における最大硬度Hｖmaxと最小硬度Hvminの差ΔHv(=Hｖmax-Hvmin)が10以下、より好ましくはΔHvが8以下となり、材質均一性に優れる。なお、ここで炭化物の粒径を0.5μm未満と限定したのは、発明者らが伸びフランジ性や硬度にミクロ組織、特に炭化物の影響が大きいと考え、これらの関係を種々検討した結果、特に炭化物の粒径を0.5μm未満と微細炭化物に限定した場合に、炭化物の体積率と鋼板特性との間に良好な相関を見出すことができたことによる。 The hot-rolled steel sheet produced by the method of the present invention is a hot-rolled steel sheet that has been subjected to hot-rolling spheroidizing annealing, and as described above, the average aspect ratio is about 5.0 or less, and the spheroidized carbide is It is a hot rolled steel sheet. Further, the volume fraction of carbides having a particle size of less than 0.5 μm is 15% or less, more preferably 10% or less of the volume fraction of all carbides, and the stretch flangeability is excellent. Here, by reducing the volume fraction of fine carbides having a particle size of less than 0.5 μm as described above, the stretch flangeability is improved. Such fine carbides become the starting point for void generation during hole expansion processing, The generated voids are connected to induce breakage, but it is considered that the starting point of void generation can be reduced by reducing the amount of carbide. Furthermore, as shown in FIG. 1 to be described later, the steel sheet of the present invention in which the volume fraction of fine carbides having a particle size of less than 0.5 μm is reduced as described above is the difference ΔHv () between the maximum hardness Hvmax and the minimum hardness Hvmin in the plate thickness direction. = Hvmax−Hvmin) is 10 or less, more preferably ΔHv is 8 or less, and the material uniformity is excellent. The reason why the carbide particle size is limited to less than 0.5 μm here is that the inventors considered that the microstructure, particularly carbide, has a large influence on stretch flangeability and hardness, and as a result of various studies on these relationships, This is because a good correlation could be found between the volume fraction of the carbide and the steel plate characteristics when the particle size of the carbide was limited to a fine carbide of less than 0.5 μm.
鋼板における粒径0.5μm以上である炭化物の量については、本発明のC量の範囲内であれば、特に問題となることはない。 The amount of carbide having a particle size of 0.5 μm or more in the steel plate is not particularly problematic as long as it is within the range of the C amount of the present invention.
表1に示す化学成分を有する鋼A〜E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し、酸洗後、同じく表2に示す条件にて熱延鋼板焼鈍を行い、板厚5.0mmの鋼板No.1〜19を製造した。なお、熱延鋼板焼鈍は非窒化性雰囲気(Ar雰囲気)で行った。 Continuously cast slabs of steels A to E having chemical components shown in Table 1 are heated to 1250 ° C., hot-rolled under the conditions shown in Table 2, pickled, and hot-rolled steel sheets under the conditions shown in Table 2 as well. Annealing was performed to produce steel plates No. 1 to 19 having a thickness of 5.0 mm. The hot-rolled steel sheet was annealed in a non-nitriding atmosphere (Ar atmosphere).
ここで、鋼板No.2、4、6、8、10は本発明例であり、鋼板No.11〜19は比較例である。そして、炭化物の粒径と体積率、板厚方向の硬度および穴拡げ率λの測定を以下の方法で行った。なお、穴拡げ率λは、伸びフランジ性を評価するための指標とした。 Here, steel plates Nos. 2, 4, 6, 8 , and 10 are examples of the present invention, and steel plates Nos. 11 to 19 are comparative examples. And the particle size and volume ratio of carbide, the hardness in the plate thickness direction, and the hole expansion ratio λ were measured by the following method. The hole expansion rate λ was used as an index for evaluating stretch flangeability.
i)炭化物の粒径と体積率の測定および球状化の観察
鋼板の圧延方向に平行な板厚断面を研磨し、板厚の1/4の位置をピクラール液(ピクリン酸+エタノール)で腐食後、走査型電子顕微鏡により倍率3000倍でミクロ組織の観察を行った。炭化物の粒径およびその体積率は、Media Cybernetics社製の画像解析ソフト“Image Pro Plus ver.4.0”(TM)を使用して画像解析にて定量化した。すなわち、各々の炭化物の粒径は、炭化物の外周上の2点と炭化物の相当楕円(炭化物と同面積で、かつ一次及び二次モーメントが等しい楕円)の重心を通る径を2度刻みに測定して平均した値である。
i) Measurement of carbide particle size and volume ratio and observation of spheroidization Polishing the plate thickness section parallel to the rolling direction of the steel plate and corroding 1/4 position of the plate thickness with Picral solution (picric acid + ethanol) The microstructure was observed with a scanning electron microscope at a magnification of 3000 times. The particle size and volume ratio of carbides were quantified by image analysis using image analysis software “Image Pro Plus ver. 4.0” (TM) manufactured by Media Cybernetics. In other words, the particle size of each carbide is measured in increments of 2 degrees through the center of gravity of two points on the outer circumference of the carbide and an equivalent ellipse of the carbide (an ellipse having the same area as the carbide and equal primary and secondary moments). The average value.
また、視野中の全ての炭化物について各々測定視野に対する面積率を求め、これを各炭化物の体積率と見なした。そして、視野ごとに粒径が0.5μm未満の炭化物ついて、体積率の合計(累積体積率)を求め、これを全炭化物の累積体積率で除して、視野ごとの体積率を求めた。前記視野ごとの体積率を50視野で求め、これを平均して、粒径が0.5μm未満の炭化物の体積率とした。 Moreover, the area ratio with respect to each measurement visual field was calculated | required about all the carbide | carbonized_materials in a visual field, and this was considered as the volume ratio of each carbide | carbonized_material. Then, the total volume ratio (cumulative volume ratio) of the carbide having a particle size of less than 0.5 μm for each field of view was obtained, and this was divided by the cumulative volume ratio of all the carbides to determine the volume ratio for each field of view. The volume ratio for each field of view was obtained from 50 fields of view, and this was averaged to obtain the volume ratio of the carbide having a particle size of less than 0.5 μm.
さらに、上記粒径の測定において、あわせて各炭化物のアスペクト比(最長径/最短径)を求めた。そして、各炭化物について求めたアスペクト比を平均(個数平均)して、平均のアスペクト比を求めた。 Furthermore, in the measurement of the particle size, the aspect ratio (longest diameter / shortest diameter) of each carbide was determined. Then, the average aspect ratio was obtained by averaging (number average) the aspect ratios obtained for each carbide.
ii)板厚方向の硬度測定
鋼板の圧延方向に平行な板厚断面を研磨し、鋼板表面から0.1mmの位置、板厚の1/8、2/8、3/8、4/8、5/8、6/8、7/8の位置、および鋼板裏面から0.1mmの位置の計9箇所をマイクロビッカース硬度計を用いて荷重4.9N(500gf)で測定した。そして、最大硬度Hvmaxと最小硬度Hvminの差ΔHv(=Hvmax-Hvmin)により板厚方向の硬度均一性を評価し、ΔHv≦10のときに硬度均一性に優れるとした。
ii) Hardness measurement in the plate thickness direction Polishing the plate thickness section parallel to the rolling direction of the steel plate, position 0.1mm from the steel plate surface, 1/8, 2/8, 3/8, 4/8, 5 of the plate thickness A total of 9 positions of / 8, 6/8, 7/8 and 0.1 mm from the back of the steel sheet were measured with a load of 4.9 N (500 gf) using a micro Vickers hardness tester. Then, the hardness uniformity in the plate thickness direction was evaluated by the difference ΔHv (= Hvmax−Hvmin) between the maximum hardness Hvmax and the minimum hardness Hvmin, and it was determined that the hardness uniformity was excellent when ΔHv ≦ 10.
iii)穴拡げ率λの測定
鋼板を、ポンチ径10mm、ダイス径12mm(クリアランス：板厚の20%)の打抜き工具を用いて打抜いた。その後、打抜いた穴を円筒平底ポンチ(径50mmφ、肩R8mm)により押し上げて穴拡げ加工し、穴縁に板厚貫通クラックが発生した時点での穴径d(mm)を測定して、次の式(3)で定義される穴拡げ率λ(%)を計算した。
λ=100×(d-10)/10 ・・・(3)
そして、同様な試験を6回行い、平均の穴拡げ率λを求めた。
iii) Measurement of hole expansion ratio λ A steel plate was punched using a punching tool having a punch diameter of 10 mm and a die diameter of 12 mm (clearance: 20% of the plate thickness). After that, the punched hole is pushed up by a cylindrical flat bottom punch (diameter 50mmφ, shoulder R8mm) to expand the hole, and the hole diameter d (mm) at the time when the plate thickness penetration crack occurs at the hole edge is measured. The hole expansion rate λ (%) defined by Equation (3) was calculated.
λ = 100 × (d-10) / 10 (3)
Then, the same test was performed 6 times to obtain an average hole expansion rate λ.
結果を表3に示す。本発明例である鋼板No.2、4、6、8、10は、いずれも粒径が0.5μm未満の炭化物の体積率が15%以下となっており、それぞれ同じ化学成分の比較例である鋼板No.11〜19に比べ、穴拡げ率λが高く、伸びフランジ性に優れている。なお、本発明例では、いずれも炭化物の平均のアスペクト比が5.0以下であり、球状化焼鈍されて炭化物が球状化していることを確認している。本発明の鋼板において、穴拡げ率λが高い原因は、上述したように粒径が0.5μm未満の微細な炭化物は穴拡げ加工時にボイド発生の起点になり、発生したボイドが連結して破断を誘発するが、その量を体積率で15%以下に低減したことによると考えられる。 The results are shown in Table 3. Steel plate Nos. 2 , 4 , 6 , 8 , and 10 that are examples of the present invention have a volume fraction of carbide of 15% or less with a particle size of less than 0.5 μm, each being a comparative example of the same chemical composition. Compared to steel plates Nos. 11 to 19, the hole expansion ratio λ is high and the stretch flangeability is excellent. In all of the examples of the present invention, it was confirmed that the average aspect ratio of the carbide was 5.0 or less, and the carbide was spheroidized by spheroidizing annealing. In the steel sheet of the present invention, the reason why the hole expansion ratio λ is high is that, as described above, fine carbide with a particle size of less than 0.5 μm becomes a starting point of void generation during hole expansion processing, and the generated voids are connected and fractured. It is thought to be due to the fact that the amount was reduced to 15% or less by volume.
図1に、ΔHvと粒径が0.5μm未満の炭化物の体積率との関係を示す。本発明例の鋼板No.1〜10のように、粒径が0.5μm未満の炭化物の体積率を15%以下にすると、上記のように伸びフランジ性に優れることに加え、ΔHvが10以下となり、優れた板厚方向の硬度均一性が得られる。なお、このように微細炭化物が硬度均一性に影響する理由として、微細炭化物がパーライトの存在していた領域に偏る傾向があることが一因であると考えられる。 FIG. 1 shows the relationship between ΔHv and the volume fraction of carbides having a particle size of less than 0.5 μm. Like the steel plate Nos. 1 to 10 of the present invention example, when the volume ratio of the carbide having a particle size of less than 0.5 μm is 15% or less, in addition to being excellent in stretch flangeability as described above, ΔHv is 10 or less. Excellent hardness uniformity in the thickness direction can be obtained. In addition, it is considered that the reason why the fine carbide influences the hardness uniformity in this way is that the fine carbide tends to be biased to a region where pearlite was present.
冷却速度が80℃/秒以上120℃/秒未満、冷却停止温度が600℃以下、かつ巻取温度が550℃以下の条件で製造された粒径が0.5μm未満の炭化物の体積率が10%以下である本発明例の鋼板No.2、4、6、8、10は、伸びフランジ性により優れているばかりでなく、ΔHvが8以下で板厚方向の硬度均一性により優れている。 10% volume fraction of carbide with a particle size of less than 0.5μm manufactured under conditions of a cooling rate of 80 ° C / second or more and less than 120 ° C / second, a cooling stop temperature of 600 ° C or less, and a coiling temperature of 550 ° C or less The following steel plate Nos. 2, 4, 6, 8, and 10 of the present invention are not only excellent in stretch flangeability but also excellent in hardness uniformity in the plate thickness direction when ΔHv is 8 or less.
F鋼(C:0.31質量%、Si:0.18質量%、Mn:0.68質量%、P:0.012質量%、S:0.0033質量%、Sol.Al:0.025質量%、N:0.0040質量%、Ar3変態点:785℃、Ac1変態点:737℃)、
G鋼(C:0.23質量%、Si:0.18質量%、Mn:0.76質量%、P:0.016質量%、S:0.0040質量%、Sol.Al:0.025質量%、N:0.0028質量%、Cr:1.2質量%、Ar3変態点:785℃、Ac1変態点:759℃)、
H鋼(C:0.32質量%、Si:1.2質量%、Mn:1.5質量%、P:0.025質量%、S:0.010質量%、Sol.Al:0.06質量%、N:0.0070質量%、Ar3変態点:804℃、Ac1変態点:746℃)、
I鋼(C:0.35質量%、Si:0.20質量%、Mn:0.68質量%、P:0.012質量%、S:0.0038質量%、Sol.Al:0.032質量%、N:0.0033質量%、Mo:0.17質量%、Cr:0.98質量%、Ar3変態点:773℃、Ac1変態点:754℃)、および、
表1に示すE鋼を、連続鋳造してスラブとした後1230℃に加熱し、表4に示す条件にて熱間圧延および熱延鋼板焼鈍を行い、板厚4.5mmの鋼板No.20〜36を製造した。なお、熱延鋼板焼鈍は非窒化性雰囲気(H2雰囲気)で行った。また、上記Ar3変態点、Ac1変態点は前記式(1)、式(2)から求め、CrあるいはMoを含有する場合は、式(1)、式(2)にCr、Moの補正項を導入して算出した。
F steel (C: 0.31 wt%, Si: 0.18 wt%, Mn: 0.68 wt%, P: 0.012 wt%, S: 0.0033 wt%, Sol. Al: 0.025 wt%, N: 0.0040 wt%, A r3 transformation (Point: 785 ° C, Ac1 transformation point: 737 ° C)
Steel G (C: 0.23 mass%, Si: 0.18 mass%, Mn: 0.76 mass%, P: 0.016 mass%, S: 0.0040 mass%, Sol.Al: 0.025 mass%, N: 0.0028 mass%, Cr: 1.2 (Mass%, Ar3 transformation point: 785 ° C, Ac1 transformation point: 759 ° C)
H Steel (C: 0.32 wt%, Si: 1.2 wt%, Mn: 1.5 wt%, P: 0.025 wt%, S: 0.010 wt%, Sol. Al: 0.06 mass%, N: 0.0070 wt%, A r3 transformation (Point: 804 ° C, Ac1 transformation point: 746 ° C)
Steel I (C: 0.35 wt%, Si: 0.20 wt%, Mn: 0.68 wt%, P: 0.012 wt%, S: 0.0038 wt%, Sol.Al: 0.032 wt%, N: 0.0033 wt%, Mo: 0.17 Mass%, Cr: 0.98 mass%, Ar3 transformation point: 773 ° C, Ac1 transformation point: 754 ° C), and
Steel E shown in Table 1 was continuously cast into a slab and then heated to 1230 ° C., subjected to hot rolling and hot-rolled steel sheet annealing under the conditions shown in Table 4, and a steel sheet No. 20 to 4.5 mm thick 36 was produced. The hot-rolled steel sheet was annealed in a non-nitriding atmosphere (H 2 atmosphere). Further, the above-mentioned Ar3 transformation point and Ac1 transformation point are determined from the above formulas (1) and (2), and when Cr or Mo is contained, correction of Cr and Mo in formula (1) and formula (2) Calculated by introducing a term.
得られた熱延鋼板に対し、実施例1と同様の方法で、炭化物の粒径と体積率、板厚方向の硬度および穴拡げ率λの測定を行った。なお、炭化物のアスペクト比について、実施例1と同様に検討したが、本発明例では全て平均のアスペクト比は5.0以下であり、球状化焼鈍されて炭化物が球状化していることを確認している。 The obtained hot-rolled steel sheet was measured for the particle size and volume ratio of carbide, the hardness in the thickness direction, and the hole expansion ratio λ in the same manner as in Example 1. The aspect ratio of the carbide was examined in the same manner as in Example 1. In all of the examples of the present invention, the average aspect ratio was 5.0 or less, and it was confirmed that the carbide was spheroidized by spheroidizing annealing. .
結果を表5に示す。 The results are shown in Table 5.
冷却速度以外の条件を一定とした鋼板No.20〜26では、冷却速度が本発明の範囲内であるNo.22〜25の伸びフランジ性、板厚方向の硬度均一性が顕著に優れている。また鋼板No.22〜25ではこれらの特性がさらに顕著に改善され、100℃前後(鋼板No.23〜25)で最良となる。 In steel plates Nos. 20 to 26 in which conditions other than the cooling rate are constant, the stretch flangeability and hardness uniformity in the plate thickness direction of Nos. 22 to 25 whose cooling rates are within the scope of the present invention are remarkably excellent. . In steel plates No. 22 to 25, these properties are further remarkably improved, and are best at around 100 ° C. (steel plates No. 23 to 25).
また冷却速度を一定として調査した鋼板No.27〜32では、冷却停止温度、巻取温度とも本発明の範囲内である鋼板No.32の伸びフランジ性、板厚方向の硬度均一性が顕著に優れている。また、冷却停止温度が600℃以下および巻取温度が550℃以下を満足する場合(鋼板No.32)は微細炭化物の体積率が10%以下となり、さらに顕著に優れた伸びフランジ性、板厚方向の硬度均一性が得られる。 In addition, in steel plates No. 27 to 32 investigated with a constant cooling rate, both the cooling stop temperature and the coiling temperature are notable for the stretch flangeability and the hardness uniformity in the thickness direction of steel plate No. 32 , which are within the scope of the present invention. Are better. In addition, when the cooling stop temperature satisfies 600 ° C or less and the coiling temperature satisfies 550 ° C or less (steel plate No. 32), the volume fraction of fine carbides is 10% or less, and the remarkably excellent stretch flangeability and thickness Uniform hardness uniformity is obtained.
鋼組成が本発明の範囲内であるＥ〜I鋼も、基本成分以外の合金元素を添加した場合(G、I鋼)を含めて、優れた伸びフランジ性、板厚方向の硬度均一性を示す。ただし、S量が多い場合(H鋼)に比べるとF鋼、G鋼およびI鋼は穴拡げ率の絶対値がさらに顕著に優れたものとなる。 E ~ I steel whose steel composition is within the scope of the present invention also has excellent stretch flangeability and hardness uniformity in the thickness direction, including the case where alloy elements other than the basic components are added (G, I steel). Show. However, compared with the case where the amount of S is large (H steel), the F steel, G steel, and I steel are significantly more excellent in the absolute value of the hole expansion rate.
Claims (4)
Cr:3.5質量%以下、Mo:0.7質量%以下。 The method for producing a high carbon hot-rolled steel sheet according to claim 1, wherein the composition of the steel further contains at least one selected from Cr and Mo having the following content ranges in addition to the above composition.
Cr: 3.5% by mass or less, Mo: 0.7% by mass or less.
Cr:3.5質量%以下、Mo:0.7質量%以下。 The high-carbon hot-rolled steel sheet according to claim 3 , wherein the composition of the steel further contains at least one selected from Cr and Mo having the following content ranges in addition to the above composition.
Cr: 3.5% by mass or less, Mo: 0.7% by mass or less.
Priority Applications (1)
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| CN101925869B (en) * | 2008-01-31 | 2012-09-26 | 丰田自动车株式会社 | Method and system for managing charge of automated guided vehicle |
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| JP5194454B2 (en) * | 2006-01-31 | 2013-05-08 | Ｊｆｅスチール株式会社 | Steel plate excellent in fine blanking workability and manufacturing method thereof |
| JP4952236B2 (en) * | 2006-12-25 | 2012-06-13 | Ｊｆｅスチール株式会社 | High carbon hot rolled steel sheet and manufacturing method thereof |
| CN101903547B (en) * | 2007-12-19 | 2012-05-23 | 杰富意钢铁株式会社 | Steel sheets and process for manufacturing the same |
| JP5157416B2 (en) * | 2007-12-19 | 2013-03-06 | Ｊｆｅスチール株式会社 | Steel sheet and manufacturing method thereof |
| JP5157417B2 (en) * | 2007-12-19 | 2013-03-06 | Ｊｆｅスチール株式会社 | Steel sheet and manufacturing method thereof |
| KR101091294B1 (en) * | 2008-12-24 | 2011-12-07 | 주식회사 포스코 | Steel Sheet With High Strength And Elongation And Method For Manufacturing Hot-Rolled Steel Sheet, Cold-Rolled Steel Sheet, Galvanized Steel Sheet And Galvannealed Steel Sheet With High Strength And Elongation |
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| JPH05195056A (en) * | 1991-08-28 | 1993-08-03 | Kobe Steel Ltd | Production of steel sheet having high ductility and high strength |
| JPH05255799A (en) * | 1992-03-11 | 1993-10-05 | Nippon Steel Corp | Hot dip plated hot rolled high strength steel sheet excellent in workability and its manufacture |
| JP3879446B2 (en) * | 2001-06-28 | 2007-02-14 | Ｊｆｅスチール株式会社 | Method for producing high carbon hot-rolled steel sheet with excellent stretch flangeability |
| JP4380471B2 (en) * | 2003-08-28 | 2009-12-09 | Ｊｆｅスチール株式会社 | High carbon hot rolled steel sheet and manufacturing method thereof |
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| CN101925869B (en) * | 2008-01-31 | 2012-09-26 | 丰田自动车株式会社 | Method and system for managing charge of automated guided vehicle |
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