JP3879447B2 - Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability - Google Patents
Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability Download PDFInfo
- Publication number
- JP3879447B2 JP3879447B2 JP2001196995A JP2001196995A JP3879447B2 JP 3879447 B2 JP3879447 B2 JP 3879447B2 JP 2001196995 A JP2001196995 A JP 2001196995A JP 2001196995 A JP2001196995 A JP 2001196995A JP 3879447 B2 JP3879447 B2 JP 3879447B2
- Authority
- JP
- Japan
- Prior art keywords
- steel sheet
- cooling
- rolled steel
- temperature
- carbide
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Expired - Fee Related
Links
Landscapes
- Heat Treatment Of Sheet Steel (AREA)
Description
【０００１】
【発明の属する技術分野】
本発明は、Ｃを0.2〜0.7質量％含有する伸びフランジ性に優れた高炭素冷延鋼板の製造方法に関する。
【０００２】
【従来の技術】
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つに、打抜き後の成形において、穴拡げ加工（バーリング）性の向上がある。この穴拡げ加工性は、プレス成形性としては伸びフランジ性で評価されている。そのため、伸びフランジ性の優れた材料が望まれている。
【０００３】
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、特開平１１-２６９５５２号公報と特開平１１-２６９５５３号公報には、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。
【０００４】
【発明が解決しようとする課題】
これらの技術では、フェライト組織が初析フェライトからなり、炭化物を含まないため柔らかく延性に優れているが、伸びフランジ性は必ずしも良好ではない。それは、打抜き加工時に、打抜き端面の近傍で初析フェライトの部分が大きく変形するため、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。その結果、これら変形量が大きく異なる粒の粒界付近に応力が集中し、球状化組織とフェライトの界面にボイドが発生する。これがクラックに成長するため、結果的には伸びフランジ性を劣化させると考えられる。
【０００５】
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、その場合は球状化した炭化物が粗大化し、加工の際にボイド発生の起点となるとともに、加工後の熱処理段階で炭化物が溶解し難くなり、焼入強度の低下につながる。
【０００６】
最近では従来にもまして、生産性向上の観点からの加工レベルに対する要求が厳しくなっている。そのため、高炭素鋼板の穴拡げ加工についても、加工度の増加等により、打抜き端面の割れが発生しやすくなっている。従って、高炭素鋼板にも高い伸びフランジ性が要求されている。
【０００７】
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、打抜き端面の割れが発生しにくい伸びフランジ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。
【０００８】
【課題を解決するための手段】
上記課題は、次の発明により解決される。その発明は、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、1.2 秒以内に冷却を開始し、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、酸洗後、冷延率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ことを特徴とする高炭素冷延鋼板の製造方法である。
【０００９】
この発明において、さらに、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する高炭素冷延鋼板の製造方法とすることもできる。
【００１０】
また、これらの発明において、さらに、冷却停止温度600℃以下で冷却を行い、巻取温度500℃以下で巻取ることを特徴とする高炭素冷延鋼板の製造方法、あるいはさらに、炭化物を含まないフェライト粒の体積率を10％以下に制御することを特徴とする高炭素冷延鋼板の製造方法とすることもできる。
【００１１】
以上の発明においてさらに、酸洗後の熱延鋼板を焼鈍温度600℃以上Ac1変態点以下で焼鈍した後、冷間圧延することを特徴とする高炭素冷延鋼板の製造方法とすることもできる。
【００１２】
これらの発明は、高炭素鋼板の伸びフランジ性に及ぼす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。その過程で、鋼板の伸びフランジ性に影響を及ぼす因子は、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。
【００１３】
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炭化物を含まないフェライト粒の体積率を、それぞれ制御することにより、高炭素冷延鋼板の伸びフランジ性が向上することがわかった。この知見に基づき、上記の組織を制御するための製造方法を検討し、伸びフランジ性に優れた高炭素冷延鋼板の製造方法を確立した。
【００１４】
以下、本発明の構成要素について説明する。
【００１５】
Ｃ含有量： 0.2〜0.7質量％
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含有量が0.2質量％未満では、熱延後の組織において初析フェライトの生成が顕著となり、炭化物の分布が不均一となる。さらにその場合、焼入後も、機械構造用部品として十分な強度が得られない。Ｃ含有量が0.7質量％を超える場合、焼鈍後でも十分な加工性が得られない。また、その場合、熱延後の鋼板の硬度が高く脆いため取扱いに不便であり、焼入後の強度も飽和する。従って、Ｃ含有量を0.2〜0.7質量％に規定する。
【００１６】
仕上温度： (Ar3変態点-20℃)以上
熱間圧延の仕上温度が(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炭化物を含まないフェライト粒が増加し、伸びフランジ性が劣化する。そこで、(Ar3変態点-20℃)以上の仕上温度で仕上圧延する。これにより、組織の均一化を図ることができ、伸びフランジ性の向上が図れる。
【００１７】
圧延後の冷却条件： 冷却速度＞120℃/秒
本発明では、変態後のフェライト粒の体積率の低減を図るため、圧延後に急冷（冷却）を行う。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、炭化物を含まないフェライト粒が15％超となり、伸びフランジ性が劣化する。従って、圧延後の冷却の冷却速度を120℃/秒超とする。
【００１８】
なお、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もできる。この場合、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でき、加工性をより一層向上できる。
【００１９】
冷却停止温度： 650℃以下
圧延後の冷却の冷却停止温度が高い場合、巻取りまでの冷却中に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、焼鈍後に微細炭化物が得られなくなり、伸びフランジ性が劣化する。冷却停止温度が650℃より高い場合、炭化物を含まないフェライト粒が15％超となり、伸びフランジ性が劣化する。従って、圧延後の冷却の冷却停止温度を650℃以下とする。さらに、炭化物を含まないフェライト粒を10％以下とする場合は、冷却停止温度を600℃以下とする。
【００２０】
巻取温度： 600℃以下
冷却後は鋼板を巻き取るが、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、焼鈍後の炭化物が粗大化し、巻取温度が600℃を超えると伸びフランジ性が劣化する。従って、巻取温度を600℃以下とする。さらに、巻取温度を500℃以下とすることにより、炭化物の分散状態が一層均一化し、極めて優れた伸びフランジ性が得られる。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。
【００２１】
熱延鋼板の焼鈍温度： 焼鈍を行う場合600℃以上Ac1変態点以下
熱延鋼板を酸洗した後、炭化物を球状化するために焼鈍（一次焼鈍）を行うこともできる。一次焼鈍の焼鈍温度が600℃未満の場合、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、やはり、焼鈍の効果が得られない。なお、優れた伸びフランジ性を得るには、焼鈍温度を680℃以上とすることが好ましい。
【００２２】
冷延鋼板の焼鈍温度： 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℃未満では未再結晶組織が残り、伸びフランジ性が劣化する。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、伸びフランジ性が劣化する。なお、優れた伸びフランジ性を得るには、焼鈍温度を680℃以上とすることが好ましい。
【００２３】
炭化物平均粒径： 0.1μm以上かつ2.0μm未満
炭化物粒径は、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。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0.1μm未満になると、硬度の上昇に伴い延性が低下し、そのため伸びフランジ性も低下する。炭化物平均粒径の増加に伴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低下する。従って、炭化物平均粒径を0.1μm以上かつ2.0μm未満に制御する。なお、炭化物平均粒径は前述のように製造条件、特に冷却停止温度、巻取温度、および焼鈍温度により制御することができる。
【００２４】
炭化物の分散状態： 炭化物を含まないフェライト粒の体積率が15％以下
炭化物の分散状態を均一とすることにより、前述のように、穴拡げ加工の際の打抜き端面における応力集中が緩和され、ボイドの発生が抑制できる。炭化物を含まないフェライト粒を、体積率にして15％以下にすることにより、炭化物の分散状態を均一にした場合と同様の効果が得られ、伸びフランジ性が著しく向上する。従って、炭化物を含まないフェライト粒の体積率を15％以下とする。さらに、炭化物を含まないフェライト粒を、体積率にして10％以下にすることで、炭化物の分散状態を一層均一化し、極めて優れた伸びフランジ性が得られる。
【００２５】
以上の発明で、炭化物を含まないというのは、通常の金属組織観察（光学顕微鏡）では炭化物が検出されないという意味である。このようなフェライト粒は、熱延後に初析フェライトとして生成した部分であり、冷間圧延＋焼鈍後の状態でも粒内の炭化物が実質的に見られない。また、体積率15％以下であれば、機械的性質（硬度）への影響も無視できる。なお、炭化物の分散状態は前述のように製造条件、特に仕上温度、圧延後の冷却の冷却速度、冷却停止温度、および巻取温度により制御することができる。
【００２６】
【発明の実施の形態】
この発明に用いる鋼は、Ｃ含有量を0.2〜0.7質量％とする他は、金属組織が前述の炭化物平均粒径および炭化物の分散状態となるものであればよい。その他の化学成分については、特に規定せず、Mn,Si,P,S,Al,Nなどの元素が通常の範囲で含有されていても問題ない。但し、好ましくは次のようにするとよい。
【００２７】
まず、Siについては、炭化物を黒鉛化し、焼入性を阻害する傾向があるので、2％以下とするのが望ましい。Mnについては、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％以下とするのが望ましい。
【００２８】
P,Sについては、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％以下であることが望ましい。また、Alについては、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％以下とするのが望ましい。Nについても、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％以下であることが望ましい。
【００２９】
さらに、目的に応じて、通常添加される範囲でB,Cr,Cu,Ni,Mo,Ti,Nb,W,V,Zr等の各種元素を添加してもよい。これらの元素は、本発明の効果には特に影響を及ぼさない。また、製造過程でSn,Pb等の各種元素が不純物として混入する場合があるが、このような不純物も本発明の効果に特に影響を及ぼすものではない。
【００３０】
本発明の高炭素鋼の成分調製には、転炉あるいは電気炉のどちらでも使用可能である。また、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。
【００３１】
このように成分調製された高炭素鋼を、造塊−分塊圧延または連続鋳造によりスラブとする。この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1280℃以下とする。
【００３２】
なお、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。また、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。
【００３３】
冷間圧延後の焼鈍については、箱焼鈍、連続焼鈍のいずれでもよい。これは、一次焼鈍の場合、即ち冷間圧延の前の熱延鋼板に焼鈍を施す場合も同様である。その後、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。
【００３４】
このようにして得られた高炭素冷延鋼板が、優れた伸びフランジ性を有する理由は次のように考えられる。伸びフランジ性には、打抜き端面の部分の内部組織が大きく影響する。特に、炭化物を含まないフェライト粒（熱延後の初析フェライト）が多い場合、球状化組織の部分との粒界からクラックが発生することが、確認されている。
【００３５】
ミクロ組織の挙動を見ると、打抜き加工時には炭化物の界面に、応力集中によるボイドの発生が顕著となる。この応力集中は、炭化物の寸法が大きいほど、また、炭化物を含まないフェライト粒が多いほど大きくなる。穴拡げ加工の際は、これらのボイドが連結しクラックとなる。
【００３６】
このように、製造条件の制御のみならず、炭化物平均粒径、および炭化物を含まないフェライト粒の占める割合を制御することにより、応力集中を小さくし、ボイドの発生を低減することができる。
【００３７】
【実施例】
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2.3mmの鋼板を製造した。ここで、鋼板No.1〜8は製造条件が本発明範囲内の本発明例であり、鋼板No.9〜16は製造条件が本発明範囲から外れる比較例である。
【００３８】
【表１】
【表２】
これらの鋼板からサンプルを採取し、炭化物平均粒径ならびに炭化物の分散状態の測定、硬度測定、および伸びフランジ性測定を行った。それぞれの試験・測定の方法および条件について以下に示す。
【００４１】
▲１▼ 炭化物平均粒径およびその分散状態
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、0.01mm2の範囲で炭化物粒径およびその分散状態（炭化物を含まないフェライト粒の体積率）の測定を行った。
【００４２】
▲２▼ 伸びフランジ性測定
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dbを測定して、次式で定義される穴拡げ率：λ(％)を求めた。
【００４３】
λ＝100×(db-d0)/d0 (1)
以上の測定結果より得られた、炭化物平均粒径、炭化物の分散状態、および伸びフランジ性を表３に示す。ここで、伸びフランジ性は式(1)の穴拡げ率λで評価した。
【００４４】
【表３】
この表３で、鋼板No.1〜8は製造条件が本発明範囲内であり、炭化物平均粒径が0.1μm以上かつ2.0μm未満、炭化物を含まないフェライト粒の体積率が15％以下の発明例である。
【００４６】
鋼板No.9〜16は製造条件が本発明範囲を外れた比較例であり、鋼板No.9,10,11,13は炭化物を含まないフェライト粒の体積率が上限15％超であり、鋼板No.12は炭化物平均粒径が下限0.1μm未満、鋼板No. 11,13,16は炭化物平均粒径が上限2.0μm以上であり、いずれも本発明の範囲外である。また、鋼板No. 14は一次焼鈍温度が低めであり、鋼板No. 15は冷圧率が低く、やはり本発明の範囲外である。
【００４７】
この表３より、発明例1〜8は、比較例9〜16に比べて、それぞれ同じ鋼種について、穴拡げ率λが20%以上向上しており、優れた伸びフランジ性を有することがわかる。特に圧延後の冷却終了温度が600℃以下、巻取温度が500℃以下、一次焼鈍600℃以上、かつ焼鈍温度が680℃以上である鋼板No.2,4,6,8は、発明例の中でもさらに優れた伸びフランジ性を有している。
【００４８】
【発明の効果】
この発明は、伸びフランジ性の向上を図るに当たって、製造条件の制御のみならず、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、打抜き時の端面におけるボイドの発生を抑制し、穴拡げ加工におけるクラックの成長を遅くすることができる。その結果、極めて伸びフランジ性に優れた高炭素冷延鋼板が提供可能となる。このような高炭素冷延鋼板を用いることにより、ギアに代表される変速機部品等の加工において加工度を高くとることができ、その結果、製造工程を省略して低コストで部品等を製造することが可能となる。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for producing a high carbon cold-rolled steel sheet excellent in stretch flangeability containing 0.2 to 0.7% by mass of C.
[0002]
[Prior art]
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 to improve the hole expansion (burring) property in the molding after punching. This hole expansion workability is evaluated as stretch flangeability as press formability. Therefore, a material excellent in stretch flangeability is desired.
[0003]
Several techniques have been studied for improving the stretch flangeability of such a high-carbon steel sheet. For example, Japanese Patent Application Laid-Open Nos. 11-269552 and 11-269553 propose methods for producing medium and high carbon steel sheets 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.
[0004]
[Problems to be solved by the invention]
In these techniques, the ferrite structure is composed of pro-eutectoid ferrite and does not contain carbides, so it is soft and excellent in ductility, but stretch flangeability is not necessarily good. This is because the pro-eutectoid ferrite part is greatly deformed in the vicinity of the punching end face during the punching process, and therefore the deformation amount differs greatly between the pro-eutectoid ferrite and the ferrite containing the spheroidized carbide. As a result, stress concentrates in the vicinity of the grain boundaries of the grains having greatly different deformation amounts, and voids are generated at the interface between the spheroidized structure and the ferrite. Since this grows into a crack, it is considered that the stretch flangeability is deteriorated as a result.
[0005]
As a countermeasure against this, it is conceivable to soften the whole by strengthening the spheroidizing annealing. However, in that case, the spheroidized carbides become coarse and become the starting point of void generation during processing, and the carbides are difficult to dissolve in the heat treatment stage after processing, leading to a decrease in quenching strength.
[0006]
Recently, demands for processing levels from the viewpoint of productivity improvement have become stricter than before. For this reason, also in the hole expansion processing of high-carbon steel sheets, cracking of the punched end surface is likely to occur due to an increase in the degree of processing. Therefore, a high stretch steel sheet is also required to have high stretch flangeability.
[0007]
In view of such circumstances, an object of the present invention is to provide a high-carbon cold-rolled steel sheet that can be manufactured without using multi-stage annealing that requires a long time and has excellent stretch flangeability in which cracking of a punched end surface is unlikely to occur. .
[0008]
[Means for Solving the Problems]
The above problem is solved by the following invention. According to the invention, steel containing 0.2 to 0.7% by mass of C is hot-rolled at a finishing temperature (Ar 3 transformation point -20 ° C.) or higher and then cooled within 1.2 seconds, with a cooling rate of 120 ° C./sec. Cooling is performed at a supercooling stop temperature of 650 ° C or lower, then wound at a coiling temperature of 600 ° C or lower, pickled, and then cold-rolled at a cold rolling rate of 30% or higher, and an annealing temperature of 600 ° C or higher and an Ac 1 transformation point a method for producing a high-carbon cold-rolled steel sheet you characterized by annealing below.
[0009]
In the present invention, furthermore, to average carbide grain size of less than 2.0μm or 0.1 [mu] m, and a manufacturing method of high carbon cold-rolled steel sheet shall be the control means controls the volume ratio of ferrite grains not containing carbide to 15% or less You can also.
[0010]
Further, in these inventions, further, cooling at a cooling stop temperature 600 ° C. or less, a method of manufacturing a high-carbon cold-rolled steel sheet shall be the wherein the winding at a coiling temperature 500 ° C. or less, or further, include carbides free ferrite grains having a volume ratio may be a method for manufacturing high carbon cold-rolled steel sheet you and controlling than 10%.
[0011]
Further, after annealing the hot-rolled steel sheet after pickling following annealing temperature 600 ° C. or higher Ac 1 transformation point, to a method for producing a high carbon cold-rolled steel sheet shall be the wherein the cold rolling in the above invention You can also.
[0012]
These inventions were made in the course of earnest research on the influence of the microstructure on the stretch flangeability of high-carbon steel sheets. In the process, it was found that factors affecting the stretch flangeability of the steel sheet have a great influence not only on the shape and amount of carbide but also on the dispersion state of carbide.
[0013]
Furthermore, it was found that the stretch flangeability of the high carbon cold-rolled steel sheet is improved by controlling the carbide average particle size as the carbide shape and the volume fraction of ferrite grains not containing carbide as the carbide dispersion state. It was. Based on this knowledge, a manufacturing method for controlling the above structure was examined, and a manufacturing method of a high-carbon cold-rolled steel sheet excellent in stretch flangeability was established.
[0014]
Hereinafter, the components of the present invention will be described.
[0015]
C content: 0.2-0.7 mass%
C is an important element that forms carbides and imparts hardness after quenching. When the C content is less than 0.2% by mass, proeutectoid ferrite is prominently formed in the structure after hot rolling, and the distribution of carbides becomes uneven. Furthermore, in that case, sufficient strength cannot be obtained as a machine structural component even after quenching. When the C content exceeds 0.7% by mass, sufficient workability cannot be obtained even after annealing. Moreover, in that case, the hardness of the steel sheet after hot rolling is high and brittle, which is inconvenient to handle, and the strength after quenching is saturated. Therefore, the C content is specified to be 0.2 to 0.7 mass%.
[0016]
Finishing temperature: The finishing temperature of (Ar 3 transformation point -20 ° C.) or higher hot rolling is less than (Ar 3 transformation point -20 ° C.), ferrite grains not containing carbide because ferrite transformation proceeds in some increases , Stretch flangeability deteriorates. Therefore, finish rolling is performed at a finishing temperature of (Ar 3 transformation point −20 ° C.) or higher. Thereby, a structure | tissue can be made uniform and stretch flangeability can be improved.
[0017]
Cooling conditions after rolling: Cooling rate> 120 ° C./second In the present invention, rapid cooling (cooling) is performed after rolling in order to reduce the volume fraction of ferrite grains after transformation. When the cooling method is slow cooling, the degree of supercooling of austenite is small and proeutectoid ferrite is generated. When the cooling rate is 120 ° C./second or less, pro-eutectoid ferrite is prominently produced, and ferrite grains not containing carbide exceed 15%, and the stretch flangeability deteriorates. Therefore, the cooling rate of cooling after rolling is set to more than 120 ° C./second.
[0018]
In addition, after finishing rolling, cooling can be started within a time period exceeding 0.1 second and less than 1.0 second. In this case, precipitates such as ferrite crystal grains and pearlite after transformation can be further refined, and workability can be further improved.
[0019]
Cooling stop temperature: 650 ° C. or less When the cooling stop temperature after rolling is high, ferrite is generated during cooling up to winding, and the lamella spacing of pearlite becomes coarse. Therefore, fine carbide cannot be obtained after annealing, and stretch flangeability deteriorates. When the cooling stop temperature is higher than 650 ° C., ferrite grains not containing carbides exceed 15%, and stretch flangeability deteriorates. Therefore, the cooling stop temperature for cooling after rolling is set to 650 ° C. or lower. Furthermore, when the ferrite grains not containing carbides are made 10% or less, the cooling stop temperature is made 600 ° C. or less.
[0020]
Winding temperature: The steel sheet is wound after cooling below 600 ° C. The higher the winding temperature, the greater the pearlite lamella spacing. Therefore, the carbide after annealing becomes coarse, and when the coiling temperature exceeds 600 ° C., the stretch flangeability deteriorates. Accordingly, the coiling temperature is set to 600 ° C. or lower. Furthermore, by setting the coiling temperature to 500 ° C. or less, the dispersion state of the carbides becomes more uniform, and extremely excellent stretch flangeability can be obtained. 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.
[0021]
Annealing temperature of hot-rolled steel sheet: When performing annealing, after pickling the hot-rolled steel sheet at 600 ° C. or higher and below the Ac 1 transformation point, annealing (primary annealing) may be performed to spheroidize the carbide. When the annealing temperature of primary annealing is less than 600 ° C., the effect of annealing cannot be obtained. On the other hand, when the annealing temperature exceeds the Ac 1 transformation point, a part is austenitized and pearlite is generated again during cooling, so that the annealing effect cannot be obtained. In order to obtain excellent stretch flangeability, the annealing temperature is preferably 680 ° C. or higher.
[0022]
Annealing temperature of cold-rolled steel sheet: 600 ° C. or higher and Ac 1 transformation point or lower After cold rolling, annealing is performed to promote recrystallization and carbide spheroidization. If the annealing temperature is less than 600 ° C., an unrecrystallized structure remains and stretch flangeability deteriorates. On the other hand, when the annealing temperature exceeds the Ac 1 transformation point, a part is austenitized and pearlite is generated again during cooling, so that the stretch flangeability is deteriorated. In order to obtain excellent stretch flangeability, the annealing temperature is preferably 680 ° C. or higher.
[0023]
Carbide average particle size: 0.1 μm or more and less than 2.0 μm Carbide particle size 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 carbide average particle size is less than 0.1 μm, the ductility decreases with an increase in hardness, and the stretch flangeability also decreases. Workability generally improves as the average carbide particle size increases, but if it exceeds 2.0 μm, stretch flangeability deteriorates due to the generation of voids in the hole expanding process. Therefore, the carbide average particle size is controlled to be 0.1 μm or more and less than 2.0 μm. The carbide average particle size can be controlled by the production conditions, particularly the cooling stop temperature, the coiling temperature, and the annealing temperature as described above.
[0024]
Dispersion state of carbide: The volume fraction of ferrite grains not containing carbide is 15% or less. By making the dispersion state of carbide uniform, stress concentration on the punched end face during hole expansion processing is reduced as described above, and voids are formed. Can be suppressed. By making the ferrite grains not containing carbide 15% or less in volume ratio, the same effect as when the carbide is uniformly dispersed can be obtained, and the stretch flangeability is remarkably improved. Therefore, the volume fraction of ferrite grains not containing carbide is set to 15% or less. Furthermore, by making the ferrite grains not containing carbides 10% or less in volume ratio, the dispersed state of carbides can be made more uniform, and extremely excellent stretch flangeability can be obtained.
[0025]
In the above invention, the fact that no carbide is contained means that the carbide is not detected by normal metallographic observation (optical microscope). Such ferrite grains are portions generated as pro-eutectoid ferrite after hot rolling, and carbides in the grains are not substantially seen even in a state after cold rolling and annealing. If the volume ratio is 15% or less, the influence on mechanical properties (hardness) can be ignored. As described above, the dispersion state of the carbides can be controlled by the manufacturing conditions, particularly the finishing temperature, the cooling rate of cooling after rolling, the cooling stop temperature, and the winding temperature.
[0026]
DETAILED DESCRIPTION OF THE INVENTION
The steel used in the present invention may be any steel as long as the metal structure is in the above-described carbide average particle size and carbide dispersion state, except that the C content is 0.2 to 0.7% by mass. Other chemical components are not particularly defined, and there is no problem even if elements such as Mn, Si, P, S, Al, and N are contained in a normal range. However, the following is preferable.
[0027]
First, for Si, it is desirable to make it 2% or less because it tends to graphitize carbides and inhibit hardenability. As for Mn, excessive addition tends to cause a decrease in ductility, so it is desirable to make it 2% or less.
[0028]
When P and S are contained excessively, the ductility is lowered and cracks are easily generated, so both are preferably 0.03% or less. In addition, when Al is added excessively, a large amount of AlN precipitates and lowers the hardenability, so it is desirable to make it 0.08% or less. N is also preferably contained in an amount of 0.01% or less because it causes a decrease in ductility when it is excessively contained.
[0029]
Furthermore, various elements such as B, Cr, Cu, Ni, Mo, Ti, Nb, W, V, and Zr may be added within a range in which they are usually added according to the purpose. These elements do not particularly affect the effects of the present invention. In addition, various elements such as Sn and Pb may be mixed as impurities during the manufacturing process, but such impurities do not particularly affect the effects of the present invention.
[0030]
Either a converter or an electric furnace can be used for preparing the components of the high carbon steel of the present invention. Further, rough rolling may be omitted during hot rolling, and finish rolling may be performed, or direct casting rolling may be performed in which a continuously cast slab is rolled as it is or for the purpose of suppressing temperature decrease.
[0031]
The high carbon steel whose components are prepared in this way is made into a slab by ingot-bundling rolling or continuous casting. The slab is hot-rolled, and at that time, the slab heating temperature is set to 1280 ° C. or less in order to avoid deterioration of the surface state due to generation of scale.
[0032]
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.
[0033]
As for annealing after cold rolling, either box annealing or continuous annealing may be used. The same applies to the case of primary annealing, that is, the case where annealing is performed on a hot-rolled steel sheet before cold rolling. Thereafter, temper rolling is performed as necessary. Since this temper rolling does not affect the hardenability, there is no particular limitation on the conditions.
[0034]
The reason why the high carbon cold-rolled steel sheet obtained in this way has excellent stretch flangeability is considered as follows. Stretch flangeability is greatly affected by the internal structure of the punched end face. In particular, it has been confirmed that cracks are generated from the grain boundary with the spheroidized structure when there are many ferrite grains not containing carbide (pre-deposited ferrite after hot rolling).
[0035]
Looking at the behavior of the microstructure, voids due to stress concentration become prominent at the carbide interface during punching. This stress concentration increases as the size of the carbide increases and as the number of ferrite grains not containing carbide increases. During the hole expanding process, these voids are connected to form a crack.
[0036]
Thus, not only the control of the manufacturing conditions but also the average grain size of carbides and the proportion of ferrite grains not containing carbides can be controlled to reduce stress concentration and reduce the generation of voids.
[0037]
【Example】
A continuous cast slab of steel having chemical components shown in Table 1 was heated to 1250 ° C., and subjected to hot rolling, cold rolling and annealing under the conditions shown in Table 2 to produce a steel plate having a thickness of 2.3 mm. Here, steel plates Nos. 1 to 8 are examples of the present invention in which the manufacturing conditions are within the scope of the present invention, and steel plates Nos. 9 to 16 are comparative examples in which the manufacturing conditions are outside the scope of the present invention.
[0038]
[Table 1]
[Table 2]
Samples were taken from these steel plates, and the average particle size of carbides and the dispersion state of carbides, hardness, and stretch flangeability were measured. Each test and measurement method and conditions are shown below.
[0041]
(1) Carbide average particle size and its dispersion state After polishing and corrosion of the sample thickness cross section, the microstructure was photographed with a scanning electron microscope, and the carbide particle size and its dispersion state (carbide in the range of 0.01 mm 2 ) The volume ratio of ferrite grains not contained was measured.
[0042]
{Circle around (2)} A sample for measuring stretch flangeability was punched using a punching tool having a punch diameter d 0 = 10 mm and a die diameter 10.46 mm (clearance 20%), and then subjected to a hole expansion test. The hole expansion test is performed by pushing up with a cylindrical flat bottom punch (50mmφ, 5R), and the hole diameter db is measured when a plate thickness penetration crack occurs at the hole edge, and the hole expansion rate defined by the following equation : Λ (%) was determined.
[0043]
λ = 100 × (db-d 0 ) / d 0 (1)
Table 3 shows the average particle diameter of carbide, the dispersion state of carbide, and stretch flangeability obtained from the above measurement results. Here, the stretch flangeability was evaluated by the hole expansion ratio λ of the formula (1).
[0044]
[Table 3]
In Table 3, steel plates Nos. 1 to 8 have the manufacturing conditions within the scope of the present invention, the carbide average particle diameter is 0.1 μm or more and less than 2.0 μm, and the volume fraction of ferrite grains not containing carbide is 15% or less. It is an example.
[0046]
Steel plates No. 9 to 16 are comparative examples in which the production conditions are outside the scope of the present invention, and steel plate Nos. 9, 10, 11 and 13 have a volume ratio of ferrite grains not containing carbide exceeding 15% at the upper limit. No. 12 has a carbide average particle size of lower limit of less than 0.1 μm, and steel plates No. 11, 13, and 16 have an upper limit of carbide average particle size of 2.0 μm or more, both of which are outside the scope of the present invention. Steel plate No. 14 has a lower primary annealing temperature, and steel plate No. 15 has a low cold pressure ratio, which is also outside the scope of the present invention.
[0047]
From Table 3, it can be seen that Invention Examples 1 to 8 have an excellent stretch flangeability because the hole expansion ratio λ is improved by 20% or more for the same steel types as compared with Comparative Examples 9 to 16, respectively. In particular, the steel sheet Nos. 2, 4, 6, and 8 having a cooling end temperature after rolling of 600 ° C. or lower, a coiling temperature of 500 ° C. or lower, primary annealing of 600 ° C. or higher, and an annealing temperature of 680 ° C. or higher are examples of the invention. Above all, it has even better stretch flangeability.
[0048]
【The invention's effect】
In this invention, in order to improve stretch flangeability, not only the production conditions are controlled, but also the carbide particle size and the dispersion state of the carbides are controlled, thereby suppressing the generation of voids at the end face during punching and expanding the holes. Crack growth in processing can be slowed. As a result, it is possible to provide a high carbon cold-rolled steel sheet that is extremely excellent in stretch flangeability. By using such a high-carbon cold-rolled steel sheet, it is possible to increase the degree of processing in the processing of transmission parts typified by gears, and as a result, the manufacturing process is omitted and parts are manufactured at low cost. It becomes possible to do.
Claims (6)
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001196995A JP3879447B2 (en) | 2001-06-28 | 2001-06-28 | Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2001196995A JP3879447B2 (en) | 2001-06-28 | 2001-06-28 | Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2003013144A JP2003013144A (en) | 2003-01-15 |
| JP3879447B2 true JP3879447B2 (en) | 2007-02-14 |
Family
ID=19034697
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2001196995A Expired - Fee Related JP3879447B2 (en) | 2001-06-28 | 2001-06-28 | Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP3879447B2 (en) |
Families Citing this family (11)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US20090126836A1 (en) * | 2005-05-29 | 2009-05-21 | Nobusuke Kariya | High Carbon Hot Rolled Steel Sheet and method for manufacturing same |
| JP5087865B2 (en) * | 2005-06-29 | 2012-12-05 | Ｊｆｅスチール株式会社 | High carbon cold-rolled steel sheet and method for producing the same |
| CN101208441A (en) | 2005-06-29 | 2008-06-25 | 杰富意钢铁株式会社 | Cold-rolled high-carbon steel plate and process for manufacturing method therefor |
| JP4600196B2 (en) * | 2005-07-26 | 2010-12-15 | Ｊｆｅスチール株式会社 | High carbon cold-rolled steel sheet with excellent workability and manufacturing method thereof |
| JP4696853B2 (en) * | 2005-10-31 | 2011-06-08 | Ｊｆｅスチール株式会社 | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet |
| KR100722395B1 (en) | 2005-12-26 | 2007-05-28 | 주식회사 포스코 | Spheroidization steel having high forming and method making of the same |
| EP1966404B1 (en) | 2005-12-26 | 2013-09-04 | Posco | Carbon steel sheet superior in formability and manufacturing method thereof |
| JP5292698B2 (en) * | 2006-03-28 | 2013-09-18 | Ｊｆｅスチール株式会社 | Extremely soft high carbon hot-rolled steel sheet and method for producing the same |
| JP5262012B2 (en) * | 2006-08-16 | 2013-08-14 | Ｊｆｅスチール株式会社 | High carbon hot rolled steel sheet and manufacturing method thereof |
| JP5128797B2 (en) * | 2006-09-12 | 2013-01-23 | 株式会社神戸製鋼所 | Method for cooling hot-rolled steel sheet |
| KR100851805B1 (en) | 2006-12-27 | 2008-08-13 | 주식회사 포스코 | Manufacturing method of high carbon steel sheet superior in impact toughness |
-
2001
- 2001-06-28 JP JP2001196995A patent/JP3879447B2/en not_active Expired - Fee Related
Also Published As
| Publication number | Publication date |
|---|---|
| JP2003013144A (en) | 2003-01-15 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| JP5292698B2 (en) | Extremely soft high carbon hot-rolled steel sheet and method for producing the same | |
| US7879163B2 (en) | Method for manufacturing a high carbon hot-rolled steel sheet | |
| JP4650006B2 (en) | High carbon hot-rolled steel sheet excellent in ductility and stretch flangeability and method for producing the same | |
| JP5050433B2 (en) | Method for producing extremely soft high carbon hot-rolled steel sheet | |
| JP5262012B2 (en) | High carbon hot rolled steel sheet and manufacturing method thereof | |
| US11401569B2 (en) | High-strength cold-rolled steel sheet and method for manufacturing same | |
| JP4600196B2 (en) | High carbon cold-rolled steel sheet with excellent workability and manufacturing method thereof | |
| JP5358914B2 (en) | Super soft high carbon hot rolled steel sheet | |
| JP3879446B2 (en) | Method for producing high carbon hot-rolled steel sheet with excellent stretch flangeability | |
| JP3879459B2 (en) | Manufacturing method of high hardenability high carbon hot rolled steel sheet | |
| JP3879447B2 (en) | Method for producing high carbon cold-rolled steel sheet with excellent stretch flangeability | |
| JP4380471B2 (en) | High carbon hot rolled steel sheet and manufacturing method thereof | |
| JP4696853B2 (en) | Method for producing high-carbon cold-rolled steel sheet with excellent workability and high-carbon cold-rolled steel sheet | |
| JP4696753B2 (en) | Method for producing high carbon cold-rolled steel sheet excellent in punching workability and high-carbon cold-rolled steel sheet | |
| JP4380469B2 (en) | High carbon hot rolled steel sheet and manufacturing method thereof | |
| JP4622609B2 (en) | Method for producing soft high workability high carbon hot rolled steel sheet with excellent stretch flangeability | |
| JP3797165B2 (en) | High carbon steel sheet for processing with small in-plane anisotropy and method for producing the same | |
| JP4403925B2 (en) | High carbon cold-rolled steel sheet and method for producing the same | |
| JP4412094B2 (en) | High carbon cold-rolled steel sheet and method for producing the same | |
| CN114763594A (en) | Cold-rolled steel sheet and method for manufacturing cold-rolled steel sheet | |
| JP3755368B2 (en) | High carbon steel plate with excellent stretch flangeability | |
| JP2003073740A (en) | Method for manufacturing high-carbon cold rolled steel sheet with high hardenability | |
| WO2020203445A1 (en) | Medium-carbon steel sheet and method for manufacturing same | |
| JP2002003943A (en) | Method for producing steel having hyperfine martensitic structure |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20040809 |
|
| A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20060704 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20060829 |
|
| RD02 | Notification of acceptance of power of attorney |
Free format text: JAPANESE INTERMEDIATE CODE: A7422 Effective date: 20060829 |
|
| RD01 | Notification of change of attorney |
Free format text: JAPANESE INTERMEDIATE CODE: A7421 Effective date: 20060920 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20061017 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20061030 |
|
| R150 | Certificate of patent or registration of utility model |
Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20101117 Year of fee payment: 4 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20111117 Year of fee payment: 5 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20111117 Year of fee payment: 5 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20121117 Year of fee payment: 6 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20131117 Year of fee payment: 7 |
|
| LAPS | Cancellation because of no payment of annual fees |