JP3887161B2 - High burring hot rolled steel sheet with excellent low cycle fatigue strength and method for producing the same - Google Patents
High burring hot rolled steel sheet with excellent low cycle fatigue strength and method for producing the same Download PDFInfo
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【０００１】
【発明の属する技術分野】
本発明は、低サイクル疲労強度に優れる高バーリング性熱延鋼板およびその製造方法に関するものであり、特に、ロードホイールをはじめとする自動車足廻り部品等の耐久性とバーリング加工性の両立が求められる部材の素材として好適な低サイクル疲労強度に優れる高バーリング性熱延鋼板およびその製造方法に関するものである。
【０００２】
【従来の技術】
近年、自動車の燃費向上などのために軽量化を目的として、Ａｌ合金等の軽金属や高強度鋼板の自動車部材への適用が進められている。ただし、Ａｌ合金等の軽金属は比強度が高いという利点があるものの鋼に比較して著しく高価であるためその適用は特殊な用途に限られている。従ってより広い範囲で自動車の軽量化を推進するためには安価な高強度鋼板の適用が強く求められている。
【０００３】
このような高強度化の要求に対してこれまでは車体重量の１／４程度を占めるホワイトボティーやパネル類に使用される冷延鋼板の分野において強度と深絞り性を兼ね備えた鋼板や焼付け硬化性のある鋼板等の開発が進められ、車体の軽量化に寄与してきた。ところが現在、軽量化の対象は車体重量の約２０％を占める構造部材や足廻り部材にシフトしてきており、これらの部材に用いる高強度熱延鋼板の開発が急務となっている。
【０００４】
ただし、高強度化は一般的に成形性（加工性）等の材料特性を劣化させるため、材料特性を劣化させずに如何に高強度化を図るかが高強度鋼板開発のカギになる。特に構造部材や足廻り部材用鋼板に求められる特性としてはバーリング加工性、疲労耐久性および耐食性等が重要であり高強度とこれら特性を如何に高次元でバランスさせるかが重要である。
例えば、ロードホイールディスク用鋼板に求められる特性としては特に疲労耐久性が重要視されている。これは、ホイールの部材特性で最も厳しい基準で管理されているのが疲労耐久性であるためである。
【０００５】
現在、これらロードホイールディスク用熱延鋼板として４４０〜５９０ＭＰａ級の鋼板が用いられているが、これら部材用鋼板に要求される強度レベルは５９０ＭＰａ級から７８０ＭＰａ級へとさらなる高強度化へ向かいつつある。一方、高強度化の目的である薄肉化はホイールに負荷されるひずみレベルの増大をもたらし、部位によっては降伏点を超えるひずみレベルでの振幅にさらされる状況が現出されてきている。
【０００６】
これまでロードホイール等足廻り部品への高強度鋼板の適用にあたって疲労耐久性を向上させるためには降伏点以下での繰返し荷重下での疲労限を重要視してきた。しかし、上述したように最近は降伏点を超えるひずみレベルでの低サイクル疲労特性の向上が望まれるようになってきている。ところが、低サイクル疲労特性を向上させるための技術については、ほとんど見受けられないのが現状である。
【０００７】
【発明が解決しようとする課題】
そこで、本発明は、低サイクル疲労強度に優れる高バーリング性熱延鋼板およびその鋼板を安価に安定して製造できる製造方法を提供することを目的とするものである。
【０００８】
【課題を解決するための手段】
本発明者らは、現在通常に採用されている連続熱間圧延設備により工業的規模で生産されている熱延鋼板の製造プロセスを念頭において、熱延鋼板の低サイクル疲労強度の向上を達成すべく鋭意研究を重ねた。その結果、体積分率最大のミクロ組織が、ベイナイト，またはフェライトおよびベイナイトの複合組織からなり、疲労試験後に観察される転位構造のうちセル構造の面積率が５０％以下であることが低サイクル疲労強度向上に非常に有効であることを新たに見出し、本発明をなしたものである。
【０００９】
即ち、本発明の要旨は、以下の通りである。
（１）質量％にて、Ｃ：０．０１〜０．２％、Ｓｉ：０．０１〜２％、Ｍｎ：０．０５〜２．５０％、Ｐ≦０．１％、Ｓ≦０．０１％、を含み、Ａｌ≦０．２％、Ｎ：０．００１〜０．１％、０．５２Ａｌ／Ｎ≦１０を満たすようにＡｌとＮを含有し、かつＣｒ、Ｍｏ、ＶをＣｒ≦２．５％、Ｍｏ≦１％、Ｖ≦０．１％、かつ（Ｃｒ＋３．５Ｍｏ＋３９Ｖ）≧０．１を満たすように含有し、残部がＦｅ及び不可避的不純物からなる鋼であって、その体積分率最大のミクロ組織が、ベイナイト，またはフェライトおよびベイナイトの複合組織からなることを特徴とする、低サイクル疲労強度に優れる高バーリング性熱延鋼板。
【００１０】
（２）前記鋼が、さらに、質量％にて、Ｃｕ：０．２〜１．２％を含有することを特徴とする、前記（１）に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
（３）前記鋼が、さらに、質量％にて、Ｂ：０．０００２〜０．００２％を含有することを特徴とする、前記（１）または（２）に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
【００１１】
（４）前記鋼が、さらに、質量％にて、Ｎｉ：０．１〜０．６％を含有することを特徴とする、前記（１）ないし（３）のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
（５）前記鋼が、さらに、質量％にて、Ｃａ：０．０００５〜０．００２％、ＲＥＭ：０．０００５〜０．０２％の一種または二種を含有することを特徴とする、前記（１）ないし（４）のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
【００１２】
（６）前記鋼が、さらに、質量％にて、Ｎｂ：０．００１〜０．１％かつＮ−０．１５Ｎｂ≧０．０００５％を含有することを特徴とする、（１）ないし（５）のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
（７） 前記鋼が、さらに、質量％にて、Ｔｉ：０．００１〜０．１％かつＮ−０．２９Ｔｉ≧０．０００５％を含有することを特徴とする、（１）ないし（５）のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
（８） 前記鋼が、さらに、質量％にて、Ｚｒ：０．００１〜０．２％の一種または二種以上を含有することを特徴とする、（１）ないし（７）のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。
【００１３】
（９）前記（１）ないし（８）のいずれか１項に記載の成分を有する鋼片の熱間圧延に際し、Ａｒ3変態点温度以上で熱間仕上圧延を終了した後、２０℃／ｓ以上の冷却速度で冷却して、４５０℃以上６５０℃以下の温度範囲の巻取温度で巻き取り、その体積率最大のミクロ組織が、ベイナイト，またはフェライトおよびベイナイトの複合組織からなることを特徴とする、低サイクル疲労強度に優れる高バーリング性熱延鋼板の製造方法。
（１０）前記熱間圧延に際し、粗圧延終了後、高圧デスケーリングを行い、Ａｒ3変態点温度以上で熱間仕上圧延を終了することを特徴とする前記（９）記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板の製造方法にある。
【００１４】
【発明の実施の形態】
以下に、本発明に至った基礎研究結果について説明する。
まず、疲労試験後の転位構造に及ぼすＡｌ、Ｎ、Ｃｒ、Ｍｏ、Ｖの添加量の影響を調査した。そのための供試材は、次のようにして準備した。すなわち、０．０６％Ｃ−０．９％Ｓｉ−１．２％Ｍｎ−０．０１％Ｐ−０．００１％Ｓを基本成分にＡｌ、Ｎ、Ｃｒ、Ｍｏ、Ｖの添加量を変化させて成分調整し溶製した鋳片をＡｒ3 変態点温度以上のいずれかの温度で板厚が３．５ｍｍになるように熱間仕上圧延を終了して後、２０℃／ｓ以上の冷却速度で冷却して、６５０℃〜常温の温度範囲で巻き取った。
【００１５】
このようにして得られた鋼板から図１に示す形状の疲労試験片を鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より圧延方向が長辺になるように採取し疲労試験に供した。ただし、疲労試験片の表面は三山仕上の研削表面とした。疲労試験は電気油圧サーボ型疲労試験機を用い、試験方法はＡＳＴＭ Ｅ６０６−９２に準じた。なお、試験条件は図２に示すように軸方向に三角波にて完全両振り引張圧縮負荷で、全ひずみ振幅を０．２〜０．６％、ひずみ速度を４．０×１０-3／ｓｅｃとした。試験はひずみ応答および応力応答の変化を記録しながら行った。疲労試験終了後、全ひずみ振幅の条件が２≦２１００×εa／ＹＰ≦４の範囲で試験を行った試験片について図３に示すように破断部近傍１／４厚の部位から透過型電子顕微鏡試料（薄膜）を加工ひずみが導入されないように採取し、透過型電子顕微鏡にて転位構造の観察を行った。ただし、透過型電子顕微鏡による観察は２０００〜１００００倍の倍率にて結晶粒を変えて１０視野以上観察した。ここでＹＰ：降伏応力または０．２％耐力（ＭＰａ）、εa：全ひずみ振幅（％）である。
【００１６】
図４および図５に観察例を示す。いずれも全ひずみ振幅εa＝０．３％の条件である。図４は本発明範囲外、図５は本発明範囲の例である。本発明範囲外の図４が典型的なセル構造を示すのに対して、本発明範囲の図５はセル構造を示さない。ここでセル構造とは疲労現象特有な転位密度が高いセル壁（ｗａｌｌ、ｖｅｉｎ、ｄｅｂｒｉｓ）に囲まれたセルが集まった構造である。また、セルとはセル壁に四方を囲まれ完全に閉じた構造のものと定義する。
【００１７】
一方、セル構造の面積率とは１試料で観察された各視野において目視または画像処理によって得られた面積率の値を観察視野毎に足し合わせ、それを観察視野数で割ったいわゆる平均値とする。さらに、本発明において転位構造を観察する箇所としては光学顕微鏡の５００倍程度で観察できる分解能においてフェライト粒を含むミクロ組織の鋼ではフェライト粒とし、ベイナイト単相組織のようにフェライト粒を含まないミクロ組織の鋼ではベイナイト組織中のセメンタイト相以外の相であるフェライト相とする。
【００１８】
本発明者らは、これらの実験結果を詳細に検討した結果、疲労試験後に観察される転位構造と低サイクル疲労強度には図６に示すように非常に強い相関があり、転位構造のうちセル構造の面積率が５０％以下であると低サイクル疲労強度が向上することを新たに知見した。また、転位構造と０．５２Ａｌ／Ｎの値およびＣｒ＋３．５Ｍｏ＋３９Ｖの値との関係においても図７に示すように強い相関関係が認められ、０．５２Ａｌ／Ｎ≦１０かつ（Ｃｒ＋３．５Ｍｏ＋３９Ｖ）≧０．１の領域においてセル構造の面積率が５０％以下になることを新たに知見した。このメカニズムは必ずしも明らかではないが以下のように推測される。通常、軟質相であるフェライト相に繰返しひずみが集中して繰返し軟化が起こり低サイクル疲労強度が低下する。従って低サイクル疲労強度を向上させるためには軟質相であるフェライト相において繰返し軟化を抑制しなければならない。
【００１９】
本発明のごとく固溶状態のＮ、ＣおよびＣｒ，Ｍｏ，Ｖを特定範囲で含有すると、進入型固溶元素であるＮやＣとＣｒ，Ｍｏ，Ｖとがフェライト相においてペアやクラスターを形成し、繰返し荷重下での転位の交差すべりを抑制することで転位の再配列（セル構造の形成）による繰返し軟化を抑制する。さらに繰返し荷重の負荷により生成する原子空孔の作用により進入型固溶元素であるＮやＣがＣｒ，Ｍｏ，Ｖのペアやクラスターから脱出し、転位を固着するため繰返し硬化が起こることで低サイクル疲労強度が向上する。
また、熱間圧延条件等を制限することによって、フェライト相において進入型固溶元素であるＮやＣの存在状態を制御し低サイクル疲労強度に優れる鋼板を製造できることも新たに知見した。
【００２０】
本発明において低サイクル疲労強度とは繰返し降伏応力を引張強度で除した値と定義する。ここで繰返し降伏応力は以下のように求めることができる。全ひずみ振幅一定での疲労試験中のひずみ応答および応力応答の変化は図８に示すようなヒステリシスループとして模式的に表される。材料は繰返しひずみにより軟化もしくは硬化しこの変化がΔσの変化として得られる。材料のΔσの値は破断寿命（Ｎｆ）の１／２の繰返し数でほとんど飽和し安定する。従って、この繰返し数でのΔσ／２をそのひずみ振幅における応力振幅σa と定義する。このσa を各ひずみ振幅について模式的に図示したものが図９である。ここでこれらのσa をひずみに対して直線近似した直線を応力‐ひずみ曲線に外挿した交点を繰返し降伏点とする。また、この交点は材料を直線弾性体（Ｈｏｏｋｅ‘ｓ ｂｏｄｙ）と仮定したときに得られる弾性直線との交点でも差し支えない。
【００２１】
次に本発明における鋼板のミクロ組織について詳細に説明する。鋼板のミクロ組織は、疲労特性とバーリング加工性を両立させるために体積分率最大のミクロ組織をフェライト、又はフェライト及びベイナイトの複合組織とした。ただし、不可避的なパーライト、残留オーステナイト、マルテンサイトを含むことを許容するものである。ここで以下ベイナイトとは、ベイニティックフェライトおよびアシュキュラーフェライト組織も含む。なお、良好な疲労特性を確保するためには、パーライトの体積分率は５％以下が望ましい。また、良好なバーリング加工性を得るためにはベイナイトの体積分率は３０％以上が望ましく、マルテンサイトの体積分率は５％未満が望ましい。さらに、良好な延性を得るためにはベイナイトの体積分率は７０％以下が望ましい。ここで、フェライト、ベイナイト、残留オーステナイト、パーライト、マルテンサイトの体積分率とは鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より切出した試料を圧延方向断面に研磨、エッチングし、光学顕微鏡を用い２００〜５００倍の倍率で観察された板厚の１／４ｔにおけるミクロ組織の面積分率で定義される。
【００２２】
続いて、本発明の化学成分の限定理由について説明する。
Ｃは、所望のミクロ組織を得るのに必要な元素である。ただし、０．２％超含有していると加工性及び溶接性が劣化するので、０．２％以下とする。一方、０．０１％未満であると強度が低下するので０．０１％以上とする。さらに、固溶状態で存在するＣはＮと同様にＣｒ、Ｍｏ、Ｖとペアやクラスターを形成するので低サイクル疲労強度向上に有効である。本発明においては、Ｎが十分に添加されており固溶Ｃ量については特に範囲を定めない。ただし、上述の全Ｃ含有量下限値以上の範囲において効果を得るために十分な固溶Ｃ量が確保されており、その範囲は０．０００５％以上、０．００４％以下であることが望ましい。
【００２３】
Ｓｉは、所望のミクロ組織を得るのに必要であるとともに固溶強化元素として強度上昇に有効である。所望の強度を得るためには、０．０１％以上含有する必要がある。しかし、２％超含有すると加工性が劣化する。そこで、Ｓｉの含有量は０．０１％以上、２％以下とする。
Ｍｎは、固溶強化元素として強度上昇に有効である。所望の強度を得るためには、０．０５％以上必要である。一方、３％超添加するとスラブ割れを生ずるため、３％以下とする。なお、Ｍｎの上限は、２．５０％（実施例の表１中、Ｇの鋼の２．５０％に基づく）とした。
【００２４】
Ｐは、不純物であり低いほど好ましく、０．１％超含有すると加工性や溶接性に悪影響を及ぼすとともに疲労特性も低下させるので、０．１％以下とする。
Ｓは、不純物であり低いほど好ましく、多すぎると局部延性やバーリング加工性を劣化させるＡ系介在物を生成するので極力低減させるべきであるが、０．０１％以下ならば許容できる範囲である。
【００２５】
Ａｌは脱酸調製剤として使用しても良い。ただし、ＡｌはＮと結合しＡｌＮを形成するため、Ｃｒ、Ｍｏ、Ｖとペアやクラスターを形成する有効なＮ量が減少するので、その添加は製造技術上無理のない範囲で必要最小限にとどめることが望ましい。すなわち、Ａｌの添加量が０．２％超ではＣｒ、Ｍｏ、Ｖとペアやクラスターを形成する有効なＮ量を確保するためにＮを多量に添加せねばならず、製造コストやＡｌＮの析出による加工性劣化の点で不利である。従ってＡｌの添加量の上限は０．２％以下とする。また、ＡｌはＡｌ2 Ｏ3 等の非金属介在物を生成し疵や局部延性の低下を招く恐れがあるのでその添加量は０．０５％以下が望ましい。さらに、製造コストや操業効率を悪化させない範囲で鋼中にＮを容易に含有させるためにはさらには０．０２％以下が望ましい。なお、Ａｌの下限は特に定めないが、０．００１％未満では製造コストや操業効率を悪化させるため、０．００１％以上とすることが望ましい。
【００２６】
Ｎは本発明において重要な元素の一つである。本発明においては、固溶状態の進入型固溶元素であるＮやＣとＣｒ，Ｍｏ，Ｖとがフェライト相においてペアやクラスターを形成し、繰返し荷重下での転位の交差すべりを抑制することで転位の再配列（セル構造の形成）による繰返し軟化を抑制し、さらに繰返し荷重の負荷により生成する原子空孔の作用により進入型固溶元素であるＮやＣがＣｒ，Ｍｏ，Ｖのペアやクラスターから脱出し転位を固着するため繰返し硬化が起こることで低サイクル疲労強度が向上する。従って、０．００１％以上の添加が必須である。一方、溶鋼中にＮを多量に添加するためには加圧等の特別な設備および操業を必要とするのでその上限は０．１％である。また、Ｎは多すぎると降伏点伸びが発生し、加工性が劣化するのでより好ましくは、０．０１％以下である。
【００２７】
さらに、ＮはＡｌと結合してＡｌＮを形成し易い元素であるので、低サイクル疲労強度の向上に寄与する固溶Ｎを確保するために０．５２Ａｌ／Ｎ≦１０と限定する。０．５２Ａｌ／Ｎの値が１０超となると、熱間圧延後の冷却過程や巻取中、容易にＡｌＮが析出するためこれを上限とする。この値が１０以下であれば熱延後の冷却速度や巻取温度を本発明の範囲で行うことによってＡｌＮの過度の析出を避けることができる。また、０．５２Ａｌ／Ｎの値が５以下では微細なＡｌＮの析出による加工性の劣化が改善されるので、より望ましくは、０．５２Ａｌ／Ｎ≦５である。
【００２８】
一方、固溶Ｎ量は上述の全Ｎ含有量範囲で調整しても良いが、固溶Ｎ量としては０．０００５〜０．００４％が望ましい。固溶Ｎが０．０００５％未満では優れた低サイクル疲労強度を得ることができず、０．００４％超では降伏点伸びが発生し加工性が劣化する。さらに、腰折れ疵発生抑制の観点から固溶Ｎ量は、０．００１２〜０．００３％が望ましい。
【００２９】
ここで固溶ＮとはＦｅ中に単独で存在するＮだけでなく、Ｃｒ，Ｍｏ、Ｖ、Ｍｎ、Ｓｉ，Ｐなどの置換型固溶元素とペアやクラスターを形成するＮも含む。固溶Ｎ量は、水素気流中加熱抽出法によって求める。この方法は試料を２００〜５００℃程度の温度域に加熱し、固溶Ｎと水素とを反応させてアンモニアとし、これを質量分析し、その分析値を換算して固溶Ｎ量を求めるものである。
また、固溶Ｎ量は、全Ｎ量からＡｌＮ、ＮｂＮ、ＶＮ、ＴｉＮ、ＢＮなどの化合物として存在するＮ量（抽出残査の化学分析から定量）を差し引いた値から求めることもできる。さらには、内部摩擦法やＦＩＭ（Ｆｉｅｌｄ Ｉｏｎ Ｍｉｃｒｏｓｃｏｐｙ)によって求めても良い。
【００３０】
Ｃｒ，Ｍｏ，Ｖは、本発明において重要な元素である。Ｃｒ，Ｍｏ，Ｖの添加量の上限は、加工性の確保とコストの点から決定され、それぞれ２．５、１、０．１％である。特にＶは添加量が多すぎると熱間圧延条件によっては窒化物を形成し、低サイクル疲労強度の向上に効果のある固溶Ｎの確保が困難となる可能性があるので０．０４％以下とするのが望ましい。一方、優れた低サイクル疲労強度を得るためには（Ｃｒ＋３．５Ｍｏ＋３９Ｖ）≧０．１を満たす必要がある。さらに、降伏点伸びの発生による加工性の劣化を回避するためには（Ｃｒ＋３．５Ｍｏ＋３９Ｖ）≧０．４がより望ましい範囲である。また、降伏点伸びの発生による加工性の劣化を回避するためには、Ｃｒ，Ｍｏ，Ｖを組み合わせて添加することがより一層効果的である。
【００３１】
Ｃｕは、固溶状態で疲労特性を改善する効果があるので必要に応じ添加する。ただし、０．２％未満では、その効果は少なく、１．２％を超えて含有すると巻取り中に析出して加工性を著しく劣化させる恐れがある。そこで、Ｃｕの含有量は０．２〜１．２％の範囲とする。
Ｂは、Ｃｕと複合添加されることによって疲労限を上昇させる効果があるので必要に応じ添加する。ただし、０．０００２％未満ではその効果を得るために不十分であり、０．００２％超添加するとスラブ割れが起こる。よって、Ｂの添加は、０．０００２％以上、０．００２％以下とする。また、Ｂを０．０００４％超添加するとＢＮが形成されるためＣｒ、Ｍｏ、Ｖとペアやクラスターを形成する有効な固溶Ｎ量が減少する可能性がある。従ってＢの添加は、０．０００２％以上０．０００４％以下がより望ましい範囲である。
【００３２】
Ｎｉは、Ｃｕ含有による熱間脆性防止のために必要に応じ添加する。ただし、０．１％未満ではその効果が少なく、０．６％を超えて添加してもその効果が飽和するので、０．１〜０．６％とする。
ＣａおよびＲＥＭは、破壊の起点となったり、加工性を劣化させる非金属介在物の形態を変化させて無害化する元素である。ただし、０．０００５％未満添加してもその効果がなく、Ｃａならば０．００２％超、ＲＥＭならば０．０２％超添加してもその効果が飽和するのでＣａ：０．０００５〜０．００２％、ＲＥＭ：０．０００５〜０．０２％添加することが望ましい。
【００３３】
Ｎｂは組織の微細化と均一化による加工性の向上や高強度化に有効であるので必要に応じて添加する。しかし、その添加量が０．００１％未満では効果を発現せず、０．１％超添加しても効果が飽和する。また、Ｎ−０．１５Ｎｂの値が０．０００５％超であると低サイクル疲労強度向上に有効な固溶Ｎの確保が困難となる。従って、Ｎｂの添加量は０．００１〜０．１％かつＮ−０．１５Ｎｂ≧０．０００５％とする。一方、Ｎｂを０．０１２％超添加するとＮｂＮを形成し易くなり、低サイクル疲労強度向上に有効な固溶Ｎの確保が困難となる恐れがあるので、０．００１〜０．０１２％がより望ましい。
【００３４】
ＴｉもＮｂと同様の効果を有するので必要に応じて添加する。しかしその添加量が０．００１％未満では効果を発現せず、０．１％超添加してもその効果は飽和する。また、Ｎ−０．２９Ｔｉの値が０．０００５％超である低サイクル疲労強度向上に有効な固溶Ｎの確保が困難となる。従って、Ｔｉの添加量は０．００１％〜０．１％かつＮ−０．２９Ｔｉ≧０．０００５％とする。一方、Ｔｉを０．０１２％超添加するとＴｉＮとして析出または晶出する可能性があり、低サイクル疲労強度向上に有効な固溶Ｎの確保が困難となる恐れがあるので、０．００１〜０．０１２％がより望ましい。
【００３５】
さらに、強度を付与するために、析出強化もしくは固溶強化元素としてＺｒを添加しても良い。ただし、０．００１％未満ではその効果を得ることができない。また、０．２％を超え添加してもその効果は飽和する。従って、Ｚｒは０．００１％〜０．２％の範囲で添加する。ただし、ＺｒはＺｒＮを形成し低サイクル疲労強度向上に有効な固溶Ｎ量を減少させる可能性があるため、０．０１％以下とすることが望ましい。
これらを主成分とする鋼にＳｎ、Ｃｏ、Ｚｎ、Ｗ、Ｍｇを合計で１％以下含有しても構わない。しかしながらＳｎは熱間圧延時に疵が発生する恐れがあるので０．０５％以下が望ましい。
【００３６】
次に、本発明の製造方法の限定理由について、以下に詳細に述べる。
本発明では、目的の成分含有量になるように成分調整した溶鋼を鋳込むことによって得たスラブを、高温鋳片のまま熱間圧延機に直送してもよいし、室温まで冷却後に加熱炉にて再加熱した後に熱間圧延してもよい。再加熱温度については特に制限はないが、１４００℃以上であると、スケールオフ量が多量になり歩留まりが低下するので、再加熱温度は１４００℃未満が望ましい。また、１０００℃未満の加熱はスケジュール上操業効率を著しく損なうため、再加熱温度は１０００℃以上が望ましい。さらに、固溶Ｎを確保するためにＡｌＮを溶解させる必要のある場合には、１１５０℃以上とすることが望ましい。
【００３７】
熱間圧延工程は、粗圧延を終了後、仕上げ圧延を行うが、最終パス温度（ＦＴ）がＡｒ3 変態点温度以上の温度域で終了する必要がある。これは、熱間圧延中に圧延温度がＡｒ3 変態点温度を切るとひずみが残留して延性が低下してしまい加工性が劣化するためである。一方、仕上げ温度の上限については特に上限を設けないがＡｒ3 変態点温度＋２００℃超では仕上げ圧延前の温度確保が事実上不可能であるので仕上げ温度の上限はＡｒ3 変態点温度＋２００℃以下が望ましい。ここで、粗圧延終了後に高圧デスケーリングを行う場合は、鋼板表面での高圧水の衝突圧Ｐ（ＭＰａ）×流量Ｌ（リットル／ｃｍ2 ）≧０．００２５の条件を満たすことが望ましい。
【００３８】
鋼板表面での高圧水の衝突圧Ｐは以下のように記述される。（「鉄と鋼」１９９１ ｖｏｌ．７７ Ｎｏ．９ ｐ１４５０参照）
Ｐ（ＭＰａ）＝５．６４×Ｐ0 ×Ｖ／Ｈ2
ただし、
Ｐ0 （ＭＰａ）：液圧力
Ｖ（リットル／ｍｉｎ）：ノズル流液量
Ｈ（ｃｍ）：鋼板表面とノズル間の距離
【００３９】
流量Ｌは以下のように記述される。
Ｌ（リットル／ｃｍ2 ）＝Ｖ／（Ｗ×ｖ）
ただし、
Ｖ（リットル／ｍｉｎ）：ノズル流液量
Ｗ（ｃｍ）：ノズル当たり噴射液が鋼板表面に当たっている幅
ｖ（ｃｍ／ｍｉｎ）：通板速度
衝突圧Ｐ×流量Ｌの上限は本発明の効果を得るためには特に定める必要はないが、ノズル流液量を増加させるとノズルの摩耗が激しくなる等の不都合が生じるため、０．０２以下とすることが望ましい。
【００４０】
さらに、仕上げ圧延後の鋼板の最大高さＲｙが１５μｍ（１５μｍＲｙ，ｌ２．５ｍｍ，ｌｎ１２．５ｍｍ）以下であることが望ましい。これは、例えば金属材料疲労設計便覧、日本材料学会編、８４ページに記載されている通り熱延または酸洗ままの鋼板の疲労強度は鋼板表面の最大高さＲｙと相関があることから明らかである。また、その後の仕上げ圧延はデスケーリング後に再びスケールが生成してしまうのを防ぐために５秒以内に行うのが望ましい。
【００４１】
本発明において仕上圧延を終了した後、所定の巻取温度にて巻取るまでの工程については、その間の冷却速度以外は特に定めないが、バーリング加工性をそれほど劣化させずに延性との両立を目指す場合は、Ａｒ3 変態点からＡｒ1 変態点までの温度域（フェライトとオーステナイトの二相域）で１〜２０秒間滞留させてもよい。ここでの滞留は、二相域でフェライト変態を促進させるために行うが、１秒未満では、二相域におけるフェライト変態が不十分なため、十分な延性が得られず、２０秒超では、パーライトが生成し、目的とする体積率最大のミクロ組織として、ベイナイト，またはフェライトおよびベイナイトの複合組織が得られない。また、１〜２０秒間の滞留をさせる温度域はフェライト変態を容易に促進させるためにはＡｒ1 変態点以上８００℃以下が望ましい。さらにＡｌＮの析出を抑制するという観点からは７００℃以下がより望ましい。さらに、１〜２０秒間の滞留時間は生産性を極端に低下させないためには１〜１０秒間とすることが望ましい。
【００４２】
また、これらの条件を満たすためには、仕上げ圧延終了後２０℃／ｓ以上の冷却速度で当該温度域に迅速に到達させることが必要である。冷却速度の上限は特に定めないが、冷却設備の能力上３００℃／ｓ以下が妥当な冷却速度である。さらに、あまりにもこの冷却速度が早いと冷却終了温度を制御できずオーバーシュートしてＡｒ1 変態点以下まで過冷却されてしまう可能性があり、延性改善の効果が失われるので、ここでの冷却速度は１５０℃／ｓ以下が望ましい。
【００４３】
仕上圧延を終了した後、所定の巻取温度（ＣＴ）にて巻取るまでの冷却速度は２０℃／ｓ以上とする。２０℃／ｓ未満の冷却速度では、パーライトが生成する恐れがあり、目的とする体積率最大のミクロ組織であるベイナイト、またはフェライトおよびベイナイトの複合組織が得られない。一方、巻取温度までの冷却速度の上限は特に定めることなく本発明の効果を得ることができるが、熱ひずみによる板そりが懸念されることから、３００℃／ｓ以下とすることが望ましい。
【００４４】
巻取温度が６５０℃超では、パーライトが生成して目的とする体積率最大のミクロ組織であるベイナイト，またはフェライトおよびベイナイトからなるミクロ組織が得られないため、巻取温度は、６５０℃以下と限定する。一方、巻取温度が４５０℃未満では、バーリング加工性に有害な残留オーステナイトまたはマルテンサイトが多量に生成する恐れがあり目的とする体積率最大のミクロ組織であるベイナイト，またはフェライトおよびベイナイトからなるミクロ組織が得られないため、巻取温度は、４５０℃以上と限定する。
熱間圧延工程終了後は必要に応じて酸洗し、その後インラインまたはオフラインで圧下率１０％以下のスキンパスまたは圧下率４０％程度までの冷間圧延を施しても構わない。
【００４５】
【実施例】
以下に、実施例により本発明をさらに説明する。
表１に示す化学成分を有するＡ〜Ｎの鋼は、転炉にて溶製して、連続鋳造後、表２に示す加熱温度（ＳＲＴ）で再加熱し、粗圧延後に同じく表２に示す仕上げ圧延温度（ＦＴ）で１．２〜５．４ｍｍの板厚に圧延した後、表２に示す巻取温度（ＣＴ）でそれぞれ巻き取った。なお一部については粗圧延後に衝突圧２．７ＭＰａ、流量０．００１リットル／ｃｍ2 の条件で高圧デスケーリングを行った。ただし、表中の化学組成についての表示は質量％である。
【００４６】
【表１】
このようにして得られた熱延板の引張試験は、供試材を、まず、ＪＩＳ Ｚ ２２０１記載の５号試験片に加工し、ＪＩＳ Ｚ ２２４１記載の試験方法に従って行った。さらに、バーリング加工性（穴拡げ性）については日本鉄鋼連盟規格ＪＦＳ Ｔ １００１−１９９６記載の穴拡げ試験方法に従って評価した。表２にその試験結果を示す。ここで、残留オーステナイト，フェライト、ベイナイト、パーライト及びマルテンサイトの体積分率とは鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より切出した試料を圧延方向断面に研磨、エッチングし、光学顕微鏡を用い２００〜５００倍の倍率で観察された板厚の１／４ｔにおけるミクロ組織の面積分率で定義される。
【００４８】
【表２】
次に、図１に示す形状の疲労試験片を鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より圧延方向が長辺になるように採取し低サイクル疲労試験に供した。ただし、疲労試験片の表面は三山仕上の研削表面とした。疲労試験は電気油圧サーボ型疲労試験機を用い、試験方法はＡＳＴＭ Ｅ６０６−９２に準じた。なお、試験条件は図２に示すように軸方向に三角波にて完全両振り引張圧縮負荷で、全ひずみ振幅を０．３〜０．６％、ひずみ速度を４．０×１０-3／ｓｅｃとした。試験はひずみ応答および応力応答の変化を記録しながら行った。疲労試験終了後、全ひずみ振幅の条件が２≦２１００×εa／ＹＰ≦４の範囲で試験を行った試験片について図３に示すように破断部近傍１／４厚の部位から透過型電子顕微鏡試料（薄膜）を加工ひずみが導入されないように採取し、透過型電子顕微鏡にて転位構造の観察を行った。表２中に、Ｓcellとしてセル構造の面積率を示す。ただし、透過型電子顕微鏡による観察は２０００〜１００００倍の倍率にて結晶粒を変えて１０視野以上観察した。ここでＹＰ：降伏応力または０．２％耐力（ＭＰａ）、εa：全ひずみ振幅（％）である。
【００５０】
鋼板の低サイクル疲労強度は、繰返し降伏応力を引張強度で除した値で評価した。ここで、繰返し降伏応力とは、破断寿命（Ｎｆ）の１／２の繰返し数での応力振幅σa をひずみに対して直線近似した直線を応力‐ひずみ曲線または弾性直線に外挿した交点とした。
本発明に沿うものは、鋼Ａ、Ｂ、Ｄ、Ｅ、Ｇ、Ｈ、Ｊ−１、Ｊ−６、Ｎ、Ｏの１０鋼であり、所定の量の鋼成分を含有し、その体積率最大のミクロ組織が、ベイナイト，またはフェライトおよびベイナイトの複合組織からなり、疲労試験後に観察される転位構造のうちセル構造の面積率が５０％以下であることを特徴とする、低サイクル疲労強度に優れる高バーリング性熱延鋼板が得られている。
【００５１】
上記以外の鋼は、以下の理由によって本発明の範囲外である。すなわち、鋼Ｃは、Ｃの含有量が本発明の範囲外であるので目的とするミクロ組織が得られず十分な疲労強度（ＴＳ）が得られていない。鋼Ｆは、Ｓの含有量が本発明の範囲外であるので十分な穴拡げ値（λ）が得られていない。鋼Ｉは、Ｐの含有量が本発明の範囲外であるので十分な低サイクル疲労強度（ＣＹＳ／ＴＳ）が得られていない。鋼Ｊ−２は、仕上圧延終了温度（ＦＴ）が本発明の範囲より低く、ひずみが残留して延性（Ｅｌ）が低い。鋼Ｊ−３は、巻取温度（ＣＴ）が本発明の範囲より高く、目的とするミクロ組織が得られず十分な穴拡げ値（λ）が得られていない。
【００５２】
鋼Ｊ−４は、巻取温度（ＣＴ）が本発明の範囲より低く、目的とするミクロ組織が得られず十分な穴拡げ値（λ）が得られていない。鋼Ｊ−５は、仕上圧延後の冷却速度（ＣＲ）が本発明の範囲より遅く、目的とするミクロ組織が得られず十分な穴拡げ値（λ）が得られていない。鋼Ｋは、０．５２Ａｌ／Ｎの値が本発明の範囲外であるので十分な低サイクル疲労強度（ＣＹＳ／ＴＳ）が得られていない。鋼Ｌは、Ｃｒ＋３．５Ｍｏ＋３９Ｖの値が本発明の範囲外であるので十分な低サイクル疲労強度（ＣＹＳ／ＴＳ）が得られていない。鋼Ｍは、Ｃの含有量が本発明の範囲外であるので目的とするミクロ組織が得られず十分な穴拡げ値（λ）が得られていない。
【００５３】
【発明の効果】
以上詳述したように、本発明は、低サイクル疲労強度に優れる高バーリング性熱延鋼板およびその製造方法に関するものであり、これらの熱延鋼板を用いることにより、自動車足廻り部品等の耐久性とバーリング加工性の両立が求められる部材においての重要な特性の一つである低サイクル疲労特性の大幅な改善が期待できるため、本発明は、工業的価値が高い発明であると言える。
【図面の簡単な説明】
【図１】疲労試験片の形状を説明する図である。
【図２】疲労試験荷重負荷方法を説明する図である。
【図３】透過型電子顕微鏡試料採取位置を説明する図である。
【図４】疲労試験後に観察される転位構造のうちセル構造の例を示す電子顕微鏡写真である。
【図５】疲労試験後に観察される転位構造のうちセル構造以外の例を示す電子顕微鏡写真である。
【図６】本発明に至る予備実験の結果を、疲労試験後のセル構造面積率と低サイクル疲労強度（繰返し降伏応力を引張強度で除した値）の関係で示す図である。
【図７】本発明に至る予備実験の結果を、０．５２Ａｌ／Ｎの値の範囲、Ｃｒ＋３．５Ｍｏ＋３９Ｖの値の範囲と疲労試験後のセル構造面積率の関係で示す図である。
【図８】疲労試験において１／２Ｎｆでの応力振幅σaを説明する図である。
【図９】疲労試験において繰返し降伏応力ＣＹＳを説明する図である。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a high burring hot-rolled steel sheet having excellent low cycle fatigue strength and a method for producing the same, and in particular, compatibility between durability and burring workability of automobile undercarriage parts such as road wheels is required. The present invention relates to a high burring hot rolled steel sheet excellent in low cycle fatigue strength suitable as a material for a member and a method for producing the same.
[0002]
[Prior art]
In recent years, application of light metals such as Al alloys and high-strength steel sheets to automobile members has been promoted for the purpose of reducing the weight in order to improve the fuel efficiency of automobiles. However, although light metals such as Al alloys have the advantage of high specific strength, their application is limited to special applications because they are significantly more expensive than steel. Therefore, in order to promote weight reduction of automobiles over a wider range, application of inexpensive high-strength steel sheets is strongly demanded.
[0003]
In response to such demands for high strength, steel sheets and bake hardening that have both strength and deep drawability in the field of cold rolled steel sheets used for white bodies and panels that occupy about 1/4 of the weight of the vehicle body. As a result, the development of steel plates and other materials has been promoted, contributing to the weight reduction of the car body. However, at present, the object of weight reduction has shifted to structural members and suspension members that account for about 20% of the weight of the vehicle body, and the development of high-strength hot-rolled steel sheets used for these members has become an urgent task.
[0004]
However, increasing strength generally degrades material properties such as formability (workability), so the key to developing high-strength steel sheets is how to increase strength without deteriorating material properties. In particular, burring workability, fatigue durability, corrosion resistance, and the like are important as properties required for structural members and steel plates for suspension members, and it is important how high strength is balanced with these properties at a high level.
For example, fatigue durability is particularly important as a characteristic required for a steel plate for a road wheel disk. This is because fatigue durability is managed according to the strictest standards of the member characteristics of the wheel.
[0005]
Currently, 440-590 MPa grade steel plates are used as hot-rolled steel plates for road wheel disks, but the strength level required for these member steel plates is increasing from 590 MPa grade to 780 MPa grade. . On the other hand, thinning, which is the purpose of increasing the strength, causes an increase in the strain level applied to the wheel, and depending on the part, a situation has been revealed that is exposed to amplitude at a strain level exceeding the yield point.
[0006]
In the past, in order to improve fatigue durability when applying high-strength steel plates to suspension parts such as road wheels, the fatigue limit under repeated loads below the yield point has been regarded as important. However, recently as described above has come to improve the low cycle fatigue characteristics of the level strain exceeds the yield point is desired. However, at present, there is almost no technique for improving low cycle fatigue characteristics.
[0007]
[Problems to be solved by the invention]
Therefore, an object of the present invention is to provide a high burring hot-rolled steel sheet excellent in low cycle fatigue strength and a production method capable of stably producing the steel sheet at low cost.
[0008]
[Means for Solving the Problems]
The present inventors achieve an improvement in the low cycle fatigue strength of a hot-rolled steel sheet in view of the manufacturing process of a hot-rolled steel sheet produced on an industrial scale by a continuous hot rolling facility that is currently normally employed. As much research as possible. As a result, the microstructure having the largest volume fraction is composed of bainite or a composite structure of ferrite and bainite, and the cell structure area ratio of the dislocation structure observed after the fatigue test is 50% or less. The present invention has been newly found out that it is very effective in improving the strength.
[0009]
That is, the gist of the present invention is as follows.
(1) In mass%, C: 0.01 to 0.2%, Si: 0.01 to 2%, Mn: 0.05 to 2.50% , P ≦ 0.1%, S ≦ 0. 01%, Al ≦ 0.2%, N: 0.001 to 0.1%, 0.52Al / N ≦ 10 so that Al and N are contained, and Cr, Mo, and V are changed to Cr. ≦ 2.5%, Mo ≦ 1%, V ≦ 0.1%, and (Cr + 3.5Mo + 39V) ≧ 0.1 so as to satisfy the balance, the balance being Fe and inevitable impurities, A high burring hot-rolled steel sheet excellent in low cycle fatigue strength, characterized in that the microstructure having the largest volume fraction is composed of bainite or a composite structure of ferrite and bainite.
[0010]
(2) High burring heat excellent in low cycle fatigue strength according to (1), wherein the steel further contains Cu: 0.2 to 1.2% by mass%. Rolled steel sheet.
(3) The steel further contains B: 0.0002 to 0.002% by mass%, and is excellent in low cycle fatigue strength according to (1) or (2). High burring hot rolled steel sheet.
[0011]
(4) The low steel according to any one of (1) to (3), wherein the steel further contains Ni: 0.1 to 0.6% by mass%. High burring hot-rolled steel sheet with excellent cycle fatigue strength.
(5) The steel further contains one or two of Ca: 0.0005 to 0.002% and REM: 0.0005 to 0.02% in mass%. The high burring hot-rolled steel sheet having excellent low cycle fatigue strength according to any one of (1) to (4).
[0012]
(6) The steel further contains, in mass%, Nb: 0.001 to 0.1% and N-0.15Nb ≧ 0.0005 % , (1) to (5 High burring hot-rolled steel sheet having excellent low cycle fatigue strength.
(7) The steel further contains, in mass%, Ti: 0.001 to 0.1% and N-0.29Ti ≧ 0.0005%, (1) to (5 High burring hot-rolled steel sheet having excellent low cycle fatigue strength.
(8) Any one of (1) to (7), wherein the steel further contains one or more of Zr: 0.001 to 0.2% in mass%. A high burring hot-rolled steel sheet having excellent low cycle fatigue strength as described in the item.
[0013]
( 9 ) Upon hot rolling of the steel slab having the component described in any one of (1) to ( 8 ) above, after finishing the hot finish rolling at an Ar 3 transformation point temperature or higher, 20 ° C./s Cooling at the above cooling rate, winding at a coiling temperature in the temperature range of 450 ° C. or more and 650 ° C. or less, and the microstructure having the maximum volume ratio is characterized by being composed of bainite or a composite structure of ferrite and bainite. A method for producing a high burring hot-rolled steel sheet having excellent low cycle fatigue strength.
( 10 ) In the hot rolling, the high cycle descaling is performed after the end of the rough rolling, and the hot finish rolling is finished at the Ar 3 transformation point temperature or higher, and the low cycle fatigue strength according to the above ( 9 ) is achieved. It exists in the manufacturing method of the outstanding high burring hot-rolled steel plate.
[0014]
DETAILED DESCRIPTION OF THE INVENTION
The basic research results that led to the present invention will be described below.
First, the influence of the addition amount of Al, N, Cr, Mo, V on the dislocation structure after the fatigue test was investigated. The test material for that purpose was prepared as follows. In other words, 0.06% C-0.9% Si-1.2% Mn-0.01% P-0.001% S is used as a basic component, and the addition amount of Al, N, Cr, Mo, V is changed. After finishing the hot finish rolling so that the thickness of the cast slab prepared by adjusting the components was melted at any temperature equal to or higher than the Ar 3 transformation temperature, the cooling rate was 20 ° C./s or higher. And then wound up in a temperature range of 650 ° C. to room temperature.
[0015]
A fatigue test piece having the shape shown in FIG. 1 was collected from the steel plate thus obtained so that the rolling direction was longer from the 1/4 W or 3/4 W position of the steel plate width, and was subjected to a fatigue test. However, the surface of the fatigue test piece was a three-sided ground surface. For the fatigue test, an electrohydraulic servo type fatigue tester was used, and the test method conformed to ASTM E606-92. In addition, as shown in FIG. 2, the test conditions were as follows: a full double swing tension compression load with a triangular wave in the axial direction, a total strain amplitude of 0.2 to 0.6%, and a strain rate of 4.0 × 10 −3 / sec. It was. The test was performed while recording changes in strain response and stress response. After completion of the fatigue test, the specimens tested under the condition of total strain amplitude of 2 ≦ 2100 × ε a / YP ≦ 4 are shown in FIG. A microscope sample (thin film) was collected so as not to introduce processing strain, and the dislocation structure was observed with a transmission electron microscope. However, in the observation with a transmission electron microscope, the crystal grains were changed at a magnification of 2000 to 10,000 times, and 10 or more fields were observed. Here, YP: yield stress or 0.2% yield strength (MPa), ε a : total strain amplitude (%).
[0016]
An observation example is shown in FIGS. In either case, the total strain amplitude ε a is 0.3%. 4 is outside the scope of the present invention, and FIG. 5 is an example of the scope of the present invention. FIG. 4 outside the scope of the present invention shows a typical cell structure, while FIG. 5 outside the scope of the present invention does not show the cell structure. Here, the cell structure is a structure in which cells surrounded by cell walls (wall, vein, debris) having a high dislocation density peculiar to a fatigue phenomenon are gathered. A cell is defined as a structure that is completely closed by being surrounded on all sides by a cell wall.
[0017]
On the other hand, the area ratio of the cell structure is a so-called average value obtained by adding the area ratio values obtained by visual observation or image processing in each visual field observed by one sample for each observation visual field, and dividing it by the number of observation visual fields. To do. Furthermore, in the present invention, the dislocation structure is observed at a resolution that can be observed at about 500 times that of an optical microscope. In a steel having a microstructure including ferrite grains, ferrite grains are used. In the structure steel, the ferrite phase is a phase other than the cementite phase in the bainite structure.
[0018]
As a result of examining these experimental results in detail, the present inventors have found that the dislocation structure observed after the fatigue test and the low cycle fatigue strength have a very strong correlation as shown in FIG. It was newly found that the low cycle fatigue strength is improved when the area ratio of the structure is 50% or less. Further, in the relationship between the dislocation structure and the values of 0.52Al / N and Cr + 3.5Mo + 39V, a strong correlation is recognized as shown in FIG. 7, and 0.52Al / N ≦ 10 and (Cr + 3.5Mo + 39V) ≧ It has been newly found that the area ratio of the cell structure is 50% or less in the region of 0.1. Although this mechanism is not necessarily clear, it is presumed as follows. Usually, repeated strain concentrates in the ferrite phase, which is a soft phase, and repeated softening occurs, resulting in a decrease in low cycle fatigue strength. Therefore, in order to improve the low cycle fatigue strength, it is necessary to repeatedly suppress softening in the ferrite phase which is a soft phase.
[0019]
When N, C and Cr, Mo, V in a solid solution state are contained in a specific range as in the present invention, N or C, which is an intrusion-type solid solution element, forms a pair or cluster in the ferrite phase. By suppressing cross slip of dislocations under repetitive loading, repetitive softening due to dislocation rearrangement (cell structure formation) is suppressed. Furthermore, the action of atomic vacancies generated by repeated loading causes the entry-type solid solution elements N and C to escape from the Cr, Mo, V pairs and clusters, and to fix dislocations, resulting in repeated hardening. Cycle fatigue strength is improved.
In addition, it has been newly found that by limiting the hot rolling conditions and the like, it is possible to manufacture a steel sheet having excellent low cycle fatigue strength by controlling the existence state of N and C, which are interstitial solid solution elements, in the ferrite phase.
[0020]
In the present invention, low cycle fatigue strength is defined as a value obtained by dividing repeated yield stress by tensile strength. Here, the repeated yield stress can be obtained as follows. Changes in strain response and stress response during a fatigue test with a constant total strain amplitude are schematically represented as a hysteresis loop as shown in FIG. The material softens or hardens due to repeated strain, and this change is obtained as a change in Δσ. The value of [Delta] [sigma] of the material is almost saturated and stable at a repetition number of 1/2 of the fracture life (Nf). Therefore, Δσ / 2 at the number of repetitions is defined as the stress amplitude σ a at the strain amplitude. FIG. 9 schematically shows this σ a for each strain amplitude. Here, the intersection point obtained by extrapolating the straight line obtained by approximating these σ a to the strain to the stress-strain curve is defined as the yield point. Further, this intersection point may be an intersection point with an elastic straight line obtained when the material is assumed to be a linear elastic body (Hooke's body).
[0021]
Next, the microstructure of the steel sheet in the present invention will be described in detail. The microstructure of the steel sheet was made of ferrite or a composite structure of ferrite and bainite in order to achieve both fatigue characteristics and burring workability. However, it is allowed to contain inevitable pearlite, retained austenite, and martensite. Here, the term “bainite” includes bainitic ferrite and ashular ferrite structures. In order to secure good fatigue characteristics, the pearlite volume fraction is desirably 5% or less. In order to obtain good burring workability, the volume fraction of bainite is desirably 30% or more, and the volume fraction of martensite is desirably less than 5%. Furthermore, in order to obtain good ductility, the volume fraction of bainite is desirably 70% or less. Here, the volume fraction of ferrite, bainite, retained austenite, pearlite, and martensite is the sample cut from the 1/4 W or 3/4 W position of the steel plate width in the rolling direction, etched and etched using an optical microscope It is defined as the area fraction of the microstructure at 1/4 t of the plate thickness observed at a magnification of 200 to 500 times.
[0022]
Then, the reason for limitation of the chemical component of this invention is demonstrated.
C is an element necessary for obtaining a desired microstructure. However, if the content exceeds 0.2%, workability and weldability deteriorate, so the content is made 0.2% or less. On the other hand, if it is less than 0.01%, the strength decreases, so 0.01% or more. Further, C present in a solid solution state is effective in improving low cycle fatigue strength because it forms a pair or cluster with Cr, Mo, V like N. In the present invention, N is sufficiently added, and there is no particular range for the amount of dissolved C. However, a sufficient amount of solid solution C is secured in order to obtain an effect in the range of the above-mentioned lower limit value of the total C content, and the range is preferably 0.0005% or more and 0.004% or less. .
[0023]
Si is effective for increasing the strength as a solid solution strengthening element as well as being necessary for obtaining a desired microstructure. In order to obtain a desired strength, it is necessary to contain 0.01% or more. However, if it exceeds 2%, workability deteriorates. Therefore, the Si content is set to 0.01% or more and 2% or less.
Mn is effective for increasing the strength as a solid solution strengthening element. In order to obtain a desired strength, 0.05% or more is necessary. On the other hand, if over 3% is added, slab cracking occurs, so the content is made 3% or less. The upper limit of Mn was 2.50% (based on 2.50% of G steel in Table 1 of Examples).
[0024]
P is preferably an impurity and is preferably as low as possible. If contained over 0.1%, the workability and weldability are adversely affected and the fatigue characteristics are also reduced.
S is an impurity and is preferably as low as possible. If it is too large, S-based inclusions that deteriorate local ductility and burring workability are generated, and should be reduced as much as possible. However, 0.01% or less is acceptable. .
[0025]
Al may be used as a deoxidizing preparation agent. However, since Al combines with N to form AlN, the effective amount of N that forms pairs and clusters with Cr, Mo, V decreases, so the addition is minimized to the extent that it is not unreasonable in terms of manufacturing technology. It is desirable to stay. That is, if the amount of Al added exceeds 0.2%, a large amount of N must be added in order to secure an effective amount of N for forming a pair or cluster with Cr, Mo, V, and the manufacturing cost and precipitation of AlN are required. It is disadvantageous in terms of workability deterioration due to. Therefore, the upper limit of the amount of Al is 0.2% or less. Moreover, since Al produces non-metallic inclusions such as Al 2 O 3 and may cause soot and a decrease in local ductility, its addition amount is preferably 0.05% or less. Furthermore, 0.02% or less is further desirable in order to easily contain N in the steel within a range not deteriorating the manufacturing cost and the operation efficiency. The lower limit of Al is not particularly defined. However, if it is less than 0.001%, the manufacturing cost and the operation efficiency are deteriorated.
[0026]
N is one of the important elements in the present invention. In the present invention, N and C, which are solid solution intrusion-type solid solution elements, and Cr, Mo, V form a pair or cluster in the ferrite phase, and suppress cross slip of dislocations under repeated loads. In addition, it suppresses repeated softening caused by rearrangement of dislocations (formation of cell structure), and further, N and C which are intrusion-type solid solution elements are Cr, Mo, V pairs due to the action of atomic vacancies generated by repeated loading. In addition, the low cycle fatigue strength is improved by the repeated hardening that takes place from the clusters and fixes the dislocations. Therefore, addition of 0.001% or more is essential. On the other hand, in order to add a large amount of N to the molten steel, special equipment such as pressurization and operation are required, so the upper limit is 0.1%. Further, if N is too much, elongation at the yield point occurs and the workability deteriorates, so it is more preferably 0.01% or less.
[0027]
Further, since N is an element that easily forms AlN by bonding with Al, the content is limited to 0.52 Al / N ≦ 10 in order to secure solid solution N that contributes to the improvement of low cycle fatigue strength. When the value of 0.52 Al / N exceeds 10, since AlN easily precipitates during the cooling process and winding after hot rolling, this is the upper limit. If this value is 10 or less, excessive precipitation of AlN can be avoided by performing the cooling rate and coiling temperature after hot rolling within the scope of the present invention. Further, when the value of 0.52Al / N is 5 or less, deterioration of workability due to fine AlN precipitation is improved, so 0.52Al / N ≦ 5 is more desirable.
[0028]
On the other hand, the solid solution N amount may be adjusted within the above-described total N content range, but the solid solution N amount is preferably 0.0005 to 0.004%. If the solute N is less than 0.0005%, excellent low cycle fatigue strength cannot be obtained, and if it exceeds 0.004%, yield point elongation occurs and workability deteriorates. Further, the amount of solid solution N is preferably 0.0012 to 0.003% from the viewpoint of suppressing the occurrence of hip creases.
[0029]
Here, the solid solution N includes not only N present alone in Fe, but also N forming a pair or cluster with a substitutional solid solution element such as Cr, Mo, V, Mn, Si, and P. The amount of solute N is determined by a heat extraction method in a hydrogen stream. In this method, a sample is heated to a temperature range of about 200 to 500 ° C., and solid solution N and hydrogen are reacted to form ammonia. This is mass analyzed, and the analytical value is converted to obtain the amount of solid solution N. It is.
Further, the amount of solute N can also be determined from a value obtained by subtracting the amount of N existing as a compound such as AlN, NbN, VN, TiN, BN, etc. (quantified from the chemical analysis of the extraction residue) from the total N amount. Further, it may be obtained by an internal friction method or FIM (Field Ion Microscopy).
[0030]
Cr, Mo and V are important elements in the present invention. The upper limit of the added amount of Cr, Mo, V is determined from the viewpoint of ensuring workability and cost, and is 2.5, 1, and 0.1%, respectively. In particular, if V is added in an excessive amount, a nitride is formed depending on hot rolling conditions, and it may be difficult to secure solid solution N effective in improving low cycle fatigue strength. Is desirable. On the other hand, in order to obtain excellent low cycle fatigue strength, it is necessary to satisfy (Cr + 3.5Mo + 39V) ≧ 0.1. Furthermore, (Cr + 3.5Mo + 39V) ≧ 0.4 is a more desirable range in order to avoid deterioration of workability due to the occurrence of yield point elongation. Further, in order to avoid the workability occurs due to the occurrence of yield point elongation, Cr, Mo, be added to the V combine seen is more effective.
[0031]
Since Cu has an effect of improving fatigue properties in a solid solution state, Cu is added as necessary. However, if the content is less than 0.2%, the effect is small, and if the content exceeds 1.2%, it may precipitate during winding to significantly deteriorate the workability. Therefore, the Cu content is in the range of 0.2 to 1.2%.
B is added as necessary because it has the effect of increasing the fatigue limit by being added in combination with Cu. However, if it is less than 0.0002%, it is insufficient for obtaining the effect, and if added over 0.002%, slab cracking occurs. Therefore, the addition of B is set to 0.0002% or more and 0.002% or less. Further, when B is added in excess of 0.0004%, BN is formed, so that there is a possibility that the effective amount of solid solution N that forms a pair or cluster with Cr, Mo, V decreases. Therefore, the addition of B is more preferably in the range of 0.0002% to 0.0004%.
[0032]
Ni is added as necessary to prevent hot brittleness due to Cu inclusion. However, if less than 0.1%, the effect is small, and even if added over 0.6%, the effect is saturated, so 0.1% to 0.6%.
Ca and REM are elements that are detoxified by changing the form of non-metallic inclusions that become the starting point of destruction or deteriorate workability. However, even if less than 0.0005% is added, there is no effect, and if Ca is more than 0.002%, and if REM is added more than 0.02%, the effect is saturated, so Ca: 0.0005 to 0 It is desirable to add 0.002% and REM: 0.0005 to 0.02%.
[0033]
Nb is effective in improving workability and increasing strength by refining and homogenizing the structure, and is added as necessary. However, if the addition amount is less than 0.001%, the effect is not exhibited, and the effect is saturated even if the addition amount exceeds 0.1%. On the other hand, if the value of N-0.15Nb is more than 0.0005%, it is difficult to secure solid solution N effective for improving low cycle fatigue strength. Therefore, the amount of Nb added is 0.001 to 0.1% and N-0.15Nb ≧ 0.0005%. On the other hand, if Nb is added in excess of 0.012%, NbN is likely to be formed, and it may be difficult to secure solid solution N effective in improving low cycle fatigue strength. desirable.
[0034]
Since Ti has the same effect as Nb, it is added as necessary. However, if the added amount is less than 0.001%, the effect is not exhibited, and even if added over 0.1%, the effect is saturated. In addition, it becomes difficult to secure solid solution N effective for improving the low cycle fatigue strength in which the value of N-0.29Ti is more than 0.0005%. Therefore, the addition amount of Ti is set to 0.001% to 0.1% and N−0.29Ti ≧ 0.0005%. On the other hand, if Ti is added in excess of 0.012%, TiN may precipitate or crystallize, and it may be difficult to secure solid solution N effective in improving low cycle fatigue strength. .012% is more desirable.
[0035]
Further, in order to impart strength, Zr may be added as a precipitation strengthening or solid solution strengthening element. However, if it is less than 0.001%, the effect cannot be obtained. Moreover, the effect is saturated even if it adds exceeding 0.2%. Therefore, Zr is added in the range of 0.001% to 0.2%. However, since Zr may form ZrN and reduce the amount of solute N effective in improving the low cycle fatigue strength, it is desirable to make it 0.01% or less.
The steel containing these as main components may contain Sn, Co, Zn, W, and Mg in total of 1% or less. However, Sn is preferably 0.05% or less because wrinkles may occur during hot rolling.
[0036]
Next, the reasons for limiting the production method of the present invention will be described in detail below.
In the present invention, a slab obtained by casting a molten steel whose components are adjusted so as to have a desired component content may be directly sent to a hot rolling mill as a high-temperature slab, or after being cooled to room temperature, a heating furnace It may be hot-rolled after reheating at. The reheating temperature is not particularly limited, but if it is 1400 ° C. or higher, the scale-off amount becomes large and the yield decreases, so the reheating temperature is preferably less than 1400 ° C. In addition, since heating below 1000 ° C. significantly impairs the operation efficiency on the schedule, the reheating temperature is desirably 1000 ° C. or higher. Furthermore, when it is necessary to dissolve AlN in order to ensure solid solution N, it is desirable that the temperature be 1150 ° C. or higher.
[0037]
In the hot rolling process, finish rolling is performed after finishing rough rolling, but the final pass temperature (FT) needs to be finished in a temperature range equal to or higher than the Ar 3 transformation point temperature. This is because if the rolling temperature falls below the Ar 3 transformation point temperature during hot rolling, strain remains and the ductility is lowered and the workability deteriorates. On the other hand, no upper limit is set for the upper limit of the finishing temperature. However, if the Ar 3 transformation point temperature is higher than + 200 ° C., it is virtually impossible to secure the temperature before the finish rolling, so the upper limit of the finishing temperature is Ar 3 transformation point temperature + 200 ° C. or lower. Is desirable. Here, when high-pressure descaling is performed after the end of rough rolling, it is desirable that the condition of high-pressure water collision pressure P (MPa) × flow rate L (liter / cm 2 ) ≧ 0.0025 on the steel plate surface is satisfied.
[0038]
The collision pressure P of high-pressure water on the steel sheet surface is described as follows. (Refer to "Iron and Steel" 1991 vol. 77 No. 9 p1450)
P (MPa) = 5.64 × P 0 × V / H 2
However,
P 0 (MPa): Fluid pressure V (L / min): Nozzle flow rate H (cm): Distance between steel plate surface and nozzle
The flow rate L is described as follows.
L (liter / cm 2 ) = V / (W × v)
However,
V (liter / min): Nozzle flow rate W (cm): Width of spray liquid per nozzle hitting the steel plate surface v (cm / min): Upper limit of plate speed collision pressure P × flow rate L is the effect of the present invention. Although it is not necessary to determine in particular in order to obtain it, it is desirable to make it 0.02 or less because an increase in the amount of nozzle flow causes problems such as increased wear on the nozzle.
[0040]
Furthermore, it is desirable that the maximum height Ry of the steel sheet after finish rolling is 15 μm (15 μm Ry, l2.5 mm, ln12.5 mm) or less. This is clear from the fact that the fatigue strength of a hot-rolled or pickled steel sheet correlates with the maximum height Ry of the steel sheet surface, as described in, for example, Metal Material Fatigue Design Handbook, edited by the Japan Society of Materials Science, page 84. is there. Further, the subsequent finish rolling is desirably performed within 5 seconds in order to prevent the scale from being generated again after descaling.
[0041]
In the present invention, after the finish rolling is finished, the process until winding at a predetermined winding temperature is not particularly defined except for the cooling rate during that time, but it is compatible with ductility without significantly degrading burring workability. When aiming, it may be retained for 1 to 20 seconds in the temperature range (two-phase region of ferrite and austenite) from the Ar 3 transformation point to the Ar 1 transformation point. The residence here is carried out in order to promote ferrite transformation in the two-phase region, but if it is less than 1 second, ferrite transformation in the two-phase region is insufficient, so that sufficient ductility cannot be obtained. Pearlite is generated, and bainite or a composite structure of ferrite and bainite cannot be obtained as the target microstructure having the maximum volume fraction. Further, the temperature range in which the residence is performed for 1 to 20 seconds is preferably from Ar 1 transformation point to 800 ° C. in order to facilitate the ferrite transformation. Furthermore, 700 degrees C or less is more desirable from a viewpoint of suppressing precipitation of AlN. Further, the residence time of 1 to 20 seconds is desirably 1 to 10 seconds so as not to extremely reduce the productivity.
[0042]
Moreover, in order to satisfy these conditions, it is necessary to quickly reach the temperature range at a cooling rate of 20 ° C./s or higher after the finish rolling. Although the upper limit of the cooling rate is not particularly defined, an appropriate cooling rate is 300 ° C./s or less because of the capacity of the cooling facility. Furthermore, if this cooling rate is too fast, the cooling end temperature cannot be controlled, and overshooting may result in overcooling to below the Ar 1 transformation point, and the effect of improving ductility is lost. The speed is desirably 150 ° C./s or less.
[0043]
After finishing rolling, the cooling rate until winding at a predetermined winding temperature (CT) is 20 ° C./s or more. When the cooling rate is less than 20 ° C./s, pearlite may be generated, and the target microstructure of bainite or the composite structure of ferrite and bainite, which has the maximum volume fraction, cannot be obtained. On the other hand, although the upper limit of the cooling rate up to the coiling temperature is not particularly defined, the effect of the present invention can be obtained.
[0044]
When the coiling temperature exceeds 650 ° C., bainite, which is the microstructure with the maximum volume ratio produced by pearlite, or a microstructure composed of ferrite and bainite cannot be obtained. Therefore, the coiling temperature is 650 ° C. or less. limit. On the other hand, if the coiling temperature is less than 450 ° C., a large amount of retained austenite or martensite harmful to burring workability may be generated, and the target microstructure of bainite or the microstructure composed of ferrite and bainite is the maximum microstructure. Since the structure cannot be obtained, the coiling temperature is limited to 450 ° C. or higher.
After completion of the hot rolling process, pickling may be performed as necessary, and then a skin pass with a reduction rate of 10% or less or cold rolling to a reduction rate of about 40% may be performed inline or offline.
[0045]
【Example】
The following examples further illustrate the present invention.
The steels A to N having chemical components shown in Table 1 are melted in a converter, re-heated at the heating temperature (SRT) shown in Table 2 after continuous casting, and also shown in Table 2 after rough rolling. After rolling to a plate thickness of 1.2 to 5.4 mm at the finish rolling temperature (FT), each was wound at the winding temperature (CT) shown in Table 2. In some cases, high pressure descaling was performed under conditions of a collision pressure of 2.7 MPa and a flow rate of 0.001 liter / cm 2 after rough rolling. However, the indication about the chemical composition in the table is mass%.
[0046]
[Table 1]
The tensile test of the hot-rolled sheet thus obtained was performed by first processing the specimen into a No. 5 test piece described in JIS Z 2201, and following the test method described in JIS Z 2241. Further, the burring workability (hole expandability) was evaluated according to the hole expansion test method described in Japan Iron and Steel Federation Standard JFS T 1001-1996. Table 2 shows the test results. Here, the volume fraction of retained austenite, ferrite, bainite, pearlite, and martensite is a sample cut from the 1/4 W or 3/4 W position of the steel plate width in the rolling direction, etched and etched using an optical microscope It is defined as the area fraction of the microstructure at 1/4 t of the plate thickness observed at a magnification of 200 to 500 times.
[0048]
[Table 2]
Next, a fatigue test piece having the shape shown in FIG. 1 was collected from the 1/4 W or 3/4 W position of the steel plate width so that the rolling direction was the longer side, and subjected to a low cycle fatigue test. However, the surface of the fatigue test piece was a three-sided ground surface. For the fatigue test, an electrohydraulic servo type fatigue tester was used, and the test method conformed to ASTM E606-92. As shown in FIG. 2, the test conditions were as follows: Fully swinging tension and compression load with a triangular wave in the axial direction, total strain amplitude of 0.3 to 0.6%, strain rate of 4.0 × 10 −3 / sec. It was. The test was performed while recording changes in strain response and stress response. After completion of the fatigue test, the specimens tested under the condition of total strain amplitude of 2 ≦ 2100 × ε a / YP ≦ 4 are shown in FIG. A microscope sample (thin film) was collected so as not to introduce processing strain, and the dislocation structure was observed with a transmission electron microscope. In Table 2, the area ratio of the cell structure is shown as S cell . However, in the observation with a transmission electron microscope, the crystal grains were changed at a magnification of 2000 to 10,000 times, and 10 or more fields were observed. Here, YP: yield stress or 0.2% yield strength (MPa), ε a : total strain amplitude (%).
[0050]
The low cycle fatigue strength of the steel sheet was evaluated by the value obtained by dividing the repeated yield stress by the tensile strength. Here, the cyclic yield stress is an intersection of a stress-strain curve or an elastic line extrapolated from a straight line obtained by linearly approximating the stress amplitude σ a at a half cycle number of the fracture life (Nf) with respect to the strain. did.
In accordance with the present invention, steels A, B, D, E, G, H, J-1, J-6, N, and O are 10 steels, containing a predetermined amount of steel components, and volume ratio thereof. The maximum microstructure is composed of bainite or a composite structure of ferrite and bainite, and the cell structure has an area ratio of 50% or less of the dislocation structure observed after the fatigue test. An excellent high burring hot-rolled steel sheet is obtained.
[0051]
Steels other than the above are outside the scope of the present invention for the following reasons. That is, steel C has a C content outside the range of the present invention, so that the target microstructure cannot be obtained and sufficient fatigue strength (TS) cannot be obtained. In Steel F, since the S content is outside the range of the present invention, a sufficient hole expansion value (λ) is not obtained. Steel I does not have a sufficient low cycle fatigue strength (CYS / TS) because the P content is outside the range of the present invention. Steel J-2 has a finish rolling finish temperature (FT) lower than the range of the present invention, strain remains, and ductility (El) is low. Steel J-3 has a coiling temperature (CT) higher than the range of the present invention, and the target microstructure cannot be obtained, so that a sufficient hole expansion value (λ) is not obtained.
[0052]
Steel J-4 has a coiling temperature (CT) lower than the range of the present invention, and the target microstructure cannot be obtained, so that a sufficient hole expansion value (λ) is not obtained. Steel J-5 has a cooling rate (CR) after finish rolling lower than the range of the present invention, and the desired microstructure cannot be obtained, so that a sufficient hole expansion value (λ) is not obtained. Steel K has a value of 0.52 Al / N that is outside the range of the present invention, so a sufficient low cycle fatigue strength (CYS / TS) is not obtained. Steel L does not have a sufficiently low cycle fatigue strength (CYS / TS) because the value of Cr + 3.5Mo + 39V is outside the range of the present invention. In steel M, the C content is outside the range of the present invention, so that the target microstructure cannot be obtained and a sufficient hole expansion value (λ) cannot be obtained.
[0053]
【The invention's effect】
As described above in detail, the present invention relates to a high burring hot-rolled steel sheet excellent in low cycle fatigue strength and a method for producing the same, and by using these hot-rolled steel sheets, durability of automobile undercarriage parts and the like is achieved. Therefore, the present invention can be said to be an invention with high industrial value because it can be expected to greatly improve low cycle fatigue characteristics, which is one of the important characteristics in a member that requires both burring workability.
[Brief description of the drawings]
FIG. 1 is a diagram illustrating the shape of a fatigue test piece.
FIG. 2 is a diagram illustrating a fatigue test load application method.
FIG. 3 is a diagram for explaining a transmission electron microscope sample collection position;
FIG. 4 is an electron micrograph showing an example of a cell structure among dislocation structures observed after a fatigue test.
FIG. 5 is an electron micrograph showing examples other than the cell structure among the dislocation structures observed after the fatigue test.
FIG. 6 is a diagram showing a result of a preliminary experiment leading to the present invention in relation to a cell structure area ratio after a fatigue test and a low cycle fatigue strength (a value obtained by dividing a cyclic yield stress by a tensile strength).
FIG. 7 is a diagram showing a result of a preliminary experiment leading to the present invention in relation to a range of 0.52 Al / N value, a range of Cr + 3.5Mo + 39V value, and a cell structure area ratio after a fatigue test.
FIG. 8 is a diagram for explaining a stress amplitude σ a at ½ Nf in a fatigue test.
FIG. 9 is a diagram for explaining repeated yield stress CYS in a fatigue test.
Claims (10)
Ｃ ：０．０１〜０．２％、
Ｓｉ：０．０１〜２％、
Ｍｎ：０．０５〜２．５０％、
Ｐ ≦０．１％、
Ｓ ≦０．０１％を含み、
Ａｌ≦０．２％、
Ｎ：０．００１〜０．１％、
０．５２Ａｌ／Ｎ≦１０を満たすようにＡｌとＮを含有し、かつＣｒ、Ｍｏ、Ｖを
Ｃｒ≦２．５％、
Ｍｏ≦１％、
Ｖ ≦０．１％、
かつ（Ｃｒ＋３．５Ｍｏ＋３９Ｖ）≧０．１を満たすように含有し、
残部がＦｅ及び不可避的不純物からなる鋼であって、その体積分率最大のミクロ組織が、ベイナイト，またはフェライトおよびベイナイトの複合組織からなることを特徴とする、低サイクル疲労強度に優れる高バーリング性熱延鋼板。In mass%
C: 0.01-0.2%
Si: 0.01-2%
Mn: 0.05-2.50 %
P ≦ 0.1%,
Including S ≦ 0.01%,
Al ≦ 0.2%,
N: 0.001 to 0.1%,
Al and N are contained so as to satisfy 0.52 Al / N ≦ 10, and Cr, Mo, and V are Cr ≦ 2.5%,
Mo ≦ 1%,
V ≦ 0.1%,
And (Cr + 3.5Mo + 39V) ≧ 0.1,
Steel with the balance being Fe and inevitable impurities, and the microstructure with the largest volume fraction is composed of bainite or a composite structure of ferrite and bainite, and high burring properties with excellent low cycle fatigue strength Hot rolled steel sheet.
Ｃａ：０．０００５〜０．００２％、
ＲＥＭ：０．０００５〜０．０２％
の一種または二種を含有することを特徴とする、請求項１ないし請求項４のいずれか１項に記載の低サイクル疲労強度に優れる高バーリング性熱延鋼板。The steel is further in mass%,
Ca: 0.0005 to 0.002%,
REM: 0.0005 to 0.02%
The high burring hot-rolled steel sheet having excellent low cycle fatigue strength according to any one of claims 1 to 4, characterized by containing one or two of the following.
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| JP4716332B2 (en) * | 2006-12-21 | 2011-07-06 | 株式会社神戸製鋼所 | Hot-rolled steel sheet excellent in stretch flangeability and surface properties and method for producing the same |
| KR101027285B1 (en) * | 2008-05-29 | 2011-04-06 | 주식회사 포스코 | High strength steel sheet for hot forming with excellent heat treatment property, hot formed hardening member and manufacturing methods thereof |
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