JP3887235B2 - High-strength steel sheet, high-strength hot-dip galvanized steel sheet, high-strength galvannealed steel sheet excellent in stretch flangeability and impact resistance, and manufacturing method thereof - Google Patents

Description

【００１５】
（７）前記（１）又は（２）のいずれかに記載の成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、下記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に５℃／ｓ以上の冷却速度で２５０℃以下に冷却することを特徴とする伸びフランジ性と耐衝突特性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
【００１６】
【数７】

【００１７】
（８）前記（１）又は（２）のいずれかに記載の成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、下記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に４５０〜５５０℃の範囲の温度域で５〜９０秒保持してから５℃／ｓ以上の冷却速度で２５０℃以下に冷却することを特徴とするフランジ性と耐衝突特性に優れた高強度合金化溶融亜鉛めっき鋼板の製造方法。
【００１８】
【数８】

【００１９】
（９）前記（１）又は（２）のいずれかに記載の成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、冷間圧延率４０％以上で冷間圧延処理を行なった後に、下記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に５℃／ｓ以上の冷却速度で２５０℃以下に冷却することを特徴とする伸びフランジ性と耐衝突特性に優れた高強度溶融亜鉛めっき鋼板の製造方法。
【００２０】
【数９】

【００２１】
（１０）前記（１）又は（２）のいずれかに記載の成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、冷間圧延率４０％以上で冷間圧延処理を行なった後に、下記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に４５０〜５５０℃の範囲の温度域で５〜９０秒保持してから５℃／ｓ以上の冷却速度で２５０℃以下に冷却することを特徴とする伸びフランジ性と耐衝突特性に優れた高強度合金化溶融亜鉛めっき鋼板の製造方法。
【００２２】
【数１０】

【００２３】
【発明の実施の形態】
以下、本発明を詳細に説明する。
本発明の伸びフランジ性と耐衝突特性に優れた高強度鋼板においては、その成分組成を、質量％で表して、Ｃが０．０４〜０．２２％、Ｓｉが１．０％以下、Ｍｎが３．０％以下、Ｐが０．０５％以下、Ｓが０．０１％以下、Ａｌが０．０１〜０．１％、及びＮが０．００１〜０．００５％となるようにし、残部をＦｅ及び不可避的不純物とする。
【００２４】
鋼板中のＣ成分は、マルテンサイトの生成に重要な元素であり、かつ本発明鋼で得られる高い塗装焼付け性を得るために不可欠な元素である。Ｃが０．０４％未満では、マルテンサイト面積率を５％以上得ることができず、Ｃの下限を０．０４％とする。また、Ｃが０．２２％を超える場合、著しく溶接性が劣化するため、Ｃの上限を０．２２％とする。
【００２５】
Ｓｉ成分は、延性を損ねず、鋼板の機械的強度を高めるのに有効な元素である。Ｓｉを１．０％を超えて添加する場合には、めっき密着性を損ねるため、これを上限とする。
【００２６】
Ｍｎ成分は、オーステナイト安定化元素であり、変態生成物を作り、鋼板の機械的強度を高めるのに有効な元素である。ただし、Ｍｎ含有量が３．０％を超える場合、溶製が困難になるばかりでなく加工性が劣化するため、Ｍｎの上限を３．０％とする。
【００２７】
Ｐ成分は、安価に鋼板の機械的強度を高める元素である。しかし、Ｐ含有量が０．０５％を超える場合には、加工後の脆性的な破壊が生じやすくなるため、Pの上限を０．０５％とする。
【００２８】
Ｓ成分は、鋼中で介在物中に存在する。Ｓ含有量が０．０１％を超える場合には、介在物量が増加し、伸びフランジ性が劣化するため、Ｓの上限を０．０１％とする。
【００２９】
Ａｌ成分は、脱酸に用いられる元素である。そのため、０．０１％未満では、脱酸が不十分となり、加工性が劣化するため、これを下限とする。また、Ａｌが０．１％を超えると、コストアップになるばかりで、脱酸効果は飽和する。そのため、上限を０．１％とする。
【００３０】
Ｎ成分は、Cと同様、マルテンサイトの生成や塗装焼付け性を得るために重要な元素である。しかし、脱酸元素であるＡｌが存在する場合、Ｎは窒化物として鋼中に存在し、延性を劣化させる。そのためＮ量は少ないほど良いが、０．００１％未満にするには、脱Ｎ工程が必要になり製造コストが上がるため、これを下限とする。また、Ｎ量が０．００５％を超えると、析出した窒化物が鋼板組織を不均一にするため、安定した強度が得られず、工業生産に不向きとなる。そこで、Ｎ量の上限は０．００５％とする。
【００３１】
本発明の高強度鋼板は、上述の成分の他に、Ｎｂ、Ｔｉ及びＶから選ばれる1種以上を合計で０．００８％以上０．０５％以下の条件で更に含むことができる。
Ｎｂ、Ｔｉ、及びＶは、いずれも炭窒化物形成元素であり、鋼中に析出物を形成し、鋼板の機械的強度、特に降伏点を高めるのに有効な元素である。Ｎｂ、Ｔｉ、Ｖから選ばれる一種以上の合計量が０．００８％に満たない場合、降伏点上昇の効果はほとんど得られず、下限を０．００８％とする。また、それらの合計量が０．０５％を超えると、加工性の劣化とともに、塗装焼付け硬化性の源である鋼中のC、Nと結びついて析出してしまい、高BH性が得られない。そのため、上限を０．０５％とする。
【００３２】
本発明の鋼板は、主相であるフェライト相と第二相であるマルテンサイト相から構成される。第二相であるマルテンサイトの粒径が２μｍを超えると、軟質なフェライト相と硬質なマルテンサイトとの界面においてボイドが発生しやすくなり、良好な伸びフランジ成形性が得られない。そのため、マルテンサイトの粒径は２μｍ以下とする。
【００３３】
また、同時に、マルテンサイト相の面積率が重要である。その面積率が５％以下では、マルテンサイト粒径を２μｍ以下にしても、その硬質なマルテンサイトの絶対量が少ない分だけ、マルテンサイト周辺への転位の導入量が少なく、したがって、塗装焼付け工程による降伏点上昇量すなわちＢＨ量も少なくなる。そのため、マルテンサイト面積率は５％以上とする。
【００３４】
ここで、前述のマルテンサイトの粒径は、任意の板断面の板厚の１／４の位置で１００μｍ×１００μｍの領域を走査型電子顕微鏡で組織観察した時の、マルテンサイトの結晶の切断長さの最大値である。また、マルテンサイト面積率は、上記の組織観察において、１００μｍ×１００μｍ領域においてマルテンサイト相が占有する面積率を意味している。
【００３５】
本発明の高強度鋼板は、上述の条件を満たす限りは、熱延鋼板として製造されても、或いは冷延鋼板として製造されてもよい。また、本発明の高強度鋼板はめっき鋼板として製造されてもよく、特に好ましくは、溶融亜鉛めっき鋼板又は合金化溶融亜鉛めっき鋼板として製造される。
【００３６】
次に、本発明における高強度鋼板、高強度溶融亜鉛めっき鋼板、及び、高強度合金化溶融亜鉛めっき鋼板の製造方法について詳細に説明する。
【００３７】
まず、本発明の高強度鋼板を熱延鋼板として製造するには、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、冷却速度４０℃／ｓ以上で２００℃以下に冷却する。
【００３８】
スラブは１１５０℃以上で６０分以上加熱する。加熱温度が１１５０℃に満たない場合、或いは、加熱時間が６０分未満の場合は、スラブ内に温度分布が生じ、長さ方向、幅方向で圧延温度差が生じ、寸法精度や鋼中組織を不均一にし、工業生産に向かない。
【００３９】
仕上げ熱間圧延は、仕上げ熱間圧延機内を８７０℃以上１０２０℃以下の温度範囲内で通過する必要がある。もし１０２０℃を超えると、オーステナイト結晶粒が粗大化し、仕上げ熱間圧延後のオーステナイト粒径も大きくなり、２μｍ以下のマルテンサイトが得られなくなる。また、８７０℃以下では、圧延中にフェライト相が生じるため、均一かつ微細なオーステナイト組織が得られず、そのため局所的にマルテンサイト粒径が２μｍを超えるものが生じてしまう。また、熱間圧延後の冷却速度が４０℃／ｓ未満の場合、マルテンサイト変態量を確保しがたく、また確保できた場合でも、冷却中のオーステナイト結晶粒の成長により変態生成するマルテンサイトの粒径が２μｍを超える。そのため、４０℃／ｓ以上が必要である。冷却終了温度が２００℃を超える場合、上記同様、マルテンサイト面積率が確保されない。そのため、冷却終了温度は２００℃以下とする。
【００４０】
本発明の高強度鋼板を冷延鋼板として製造するには、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、冷間圧延率４０％以上で冷間圧延処理を行ない、下記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で４００℃未満まで冷却する。
【００４１】
【数１１】

【００４２】
冷間圧延率が４０％未満では、板厚中心まで冷間歪みが均一に導入されず、焼鈍後の中心層における冷間歪み量が少ない。そのため、再結晶により生成するフェライト粒より、隣接する冷間圧延状態の結晶粒が併呑し合う異常フェライト粒が成長するため、再結晶により生成するオーステナイト粒径も不均一となる。したがって、オーステナイト粒より変態生成するマルテンサイトの粒径が十分に小さくならず、十分な伸びフランジ性が得られない。
【００４３】
冷間圧延後の焼鈍工程において、（Ｉ）式を満足するようなα相、γ相の二相域に加熱・保持することが重要である。（Ｉ）式の上限以上の温度域では、γ相単相になるため、γ／α変態により生成するフェライト相の核生成頻度が減少するため、フェライト粒径が粗大化し、同時にオーステナイト粒も成長するため、２μｍ以下のマルテンサイト粒径が得られない。また、（Ｉ）式の下限未満の温度域では、γ相単相になるような温度域における焼鈍ほどではないが、フェライト結晶粒の成長と同時にオーステナイト粒も成長するため、所望とするマルテンサイト粒径が得られない。
【００４４】
また、保持時間が３秒未満の場合、組織の不均一が生じ易く、鋼板の長手方向或いは幅方向での材質のバラツキにつながる。そのため、保持時間は３秒以上とする。二相域焼鈍後の冷却速度は２℃／ｓ未満であるとパーライト変態が生じ、マルテンサイト量が確保し難くなり、２００℃／ｓを超える場合は鋼板形状が著しく悪化するため、冷却速度は２〜２００℃／ｓとする。冷却終了温度は４００℃未満が好ましい。４００℃以上の場合、パーライト変態相が生成し、十分なマルテンサイト相面積率が得られないためである。
【００４５】
本発明の高強度鋼板を溶融亜鉛めっき鋼板として製造するには、前述したように、本発明の高強度鋼板を、（１）熱延鋼板とした後、溶融亜鉛めっき処理を施す方法、（２）冷延鋼板とした後、溶融亜鉛めっき処理を施す方法、のいずれの方法でも採用することができる。
【００４６】
前記（１）の場合は、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、次いで脱スケール処理後、上記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に５℃／ｓ以上の冷却速度で２５０℃以下に冷却する。
【００４７】
前記（２）の場合は、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、冷間圧延率４０％以上で冷間圧延処理を行なった後に、上記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に５℃／ｓ以上の冷却速度で２５０℃以下に冷却する。
【００４８】
脱スケール処理後、あるいはそれに続く冷間圧延処理後の焼鈍工程において、
（Ｉ）式を満足するような焼鈍温度に加熱、保持する理由は既に述べた通りである。
【００４９】
二相域焼鈍後の冷却速度は２〜２００℃／ｓとする。２℃／ｓ未満であるとパーライト変態が生じ、マルテンサイト量が確保しがたくなり、２００℃／ｓを超える場合、鋼板形状が著しく悪化する。
【００５０】
冷却終了温度は３５０℃〜５００℃とする。冷却終了温度が３５０℃未満であると、亜鉛浴温度が下がるため、浴温を加熱するなどコスト高となり、５００℃を超える場合は、γ相よりベイナイト変態が起こり、所望のマルテンサイト面積率が得られない。
【００５１】
めっき後の冷却工程は、冷却速度が５℃／ｓ未満に達しない場合や２５０℃以下まで冷却を施さない場合、パーライト変態相の生成が起こり、マルテンサイト相が得られない。そこで、５℃／ｓ以上で２５０℃以下まで冷却する。
ここで、溶融亜鉛めっき方法は特に限定されるものではなく、例えば、ライン内焼鈍方式の連続溶融亜鉛めっき設備を用いるなど、通常行われている方法でよい。
【００５２】
本発明の高強度鋼板を合金化溶融亜鉛めっき鋼板として製造するには、前述した溶融亜鉛めっき鋼板の製造方法と同様に、（１）熱延鋼板とした後、溶融亜鉛めっき処理を行い、更に、合金化処理を施す方法、（２）冷延鋼板として製造した後、溶融亜鉛めっき処理を行い、更に、合金化処理を施す方法、のいずれの方法でも採用することができる。
【００５３】
前記（１）の場合は、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、次いで脱スケール処理後、上記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に４５０〜５５０℃の範囲の温度域で５〜９０秒保持してから５℃／ｓ以上の冷却速度で２５０℃以下に冷却する。
【００５４】
前記（２）の場合は、本発明における成分組成を満足するスラブを鋳造凝固後、１１５０℃以上で６０分以上加熱し、仕上げ熱間圧延を８７０℃以上１０２０℃以下で行い、脱スケール処理後、冷間圧延率４０％以上で冷間圧延処理を行なった後に、上記（Ｉ）式を満足するような焼鈍温度で３秒以上焼鈍した後、２〜２００℃／ｓの冷却速度で３５０〜５００℃まで冷却し、次いで溶融亜鉛めっきを施し、その後に４５０〜５５０℃の範囲の温度領域で５〜９０秒保持してから５℃／ｓ以上の冷却速度で２５０℃以下に冷却する。
【００５５】
合金化溶融亜鉛めっき鋼板の製造は、溶融亜鉛めっき処理に続いて合金化処理を施すことを除いて、溶融亜鉛めっき鋼板の製造と同様である。
【００５６】
合金化処理においては、加熱温度４５０〜５５０℃で５〜９０秒保持する必要がある。
４５０〜５５０℃の温度域での保持が５秒未満の場合は、十分な鉄と亜鉛の合金層を形成できず、合金化溶融亜鉛めっき鋼板としての特性が得られない。また、前記温度域で９０秒を超えての保持は、フェライト粒界に鉄炭化物のセメンタイト（Ｆｅ３Ｃ）が生成するため、プレス加工により導入される転位と相互作用するC量が減少し、ＢＨ量が減少する。そのため、保持時間は５〜９０秒の範囲とする。
【００５７】
合金化溶融亜鉛めっきを施す方法は、特に限定されるものではなく、例えば、ライン内焼鈍方式の連続溶融亜鉛めっき設備を用いるなど、通常行われている方法でよい。
【００５８】
【実施例】
次に、実施例により本発明を更に説明する。言うまでもなく、本発明はこれらの実施例に限定されるものではない。
【００５９】
表１に示す成分組成を有する鋼を鋳造凝固後、そのまま、または、いったん室温まで冷却した後に表２に示す条件で熱間圧延を施し、熱延鋼板を製造した。鋼片は１１５０℃以上で６０分以上加熱した後、熱間圧延を施した。これらの熱延鋼板の一部に対し、更に脱スケール処理後、表２に示す条件で、焼鈍処理につづき溶融亜鉛めっき処理または合金化溶融亜鉛めっき処理を施した。あるいは、熱延鋼板の別の一部を脱スケール処理後、表２に示す条件で冷間圧延し、焼鈍処理を施し、冷延鋼板を製造した。こうして得られた冷延鋼板の一部に対し、表２に示す条件で、焼鈍処理につづき溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理を施した。
【００６０】
得られた各鋼板のミクロ組織を顕微鏡観察するとともに、圧延方向に垂直に試験片を採取し、引張試験により、降伏強さ（ＹＰ）、引張強さ（ＴＳ）、塗装焼付硬化（ＢＨ）量を調べた。ここで、降伏強さ及び引張り強さは、ＪＩＳ Ｚ ２２０１記載の５号試験片を用い、ＪＩＳ Ｚ ２２４１記載の方法に従って測定した。また、ＢＨ量とは、２％の引張予変形後の応力から、１７０℃で２０分の時効処理後に再引張した時の降伏応力への増加量である。また、２００ｍｍ角の試験片を切り出し、穴径２５ｍｍの穴に対し、平底円筒パンチ（直径１００ｍｍ）を用いた穴広げ試験を行い、伸びフランジ性を評価した。伸びフランジ性は、（穴縁に割れが生じた時の穴径−初期の穴径）÷初期の穴径で定義される穴広げ率λで評価し、λが５０％を超える場合を合格とした。
【００６１】
以上の結果を表３に示す。表３において、ＹＰ＋ＷＨ＋ＢＨとは、降伏強さ（ＹＰ）に２％の予歪み変形による降伏強さからの強さ増分（ＷＨ）と塗装焼付硬化量（ＢＨ）を加えた量で、塗装焼付け工程後の降伏応力（上述のＢＨ量の測定で記述した「再引張した時の降伏応力」）に相当する。したがって、この値が大きいほど、衝突する際の降伏強さが大きく、耐衝突特性が良好になる。本発明では、原板での成形性が同程度である部材のプレス・塗装焼付け後の耐衝突特性を表す指標として、ＴＳに対する（ＹＰ＋ＷＨ＋ＢＨ）の比、即ち、（ＹＰ＋ＷＨ＋ＢＨ）／ＴＳを定義した。従来の鋼板では、高くても８５％程度であり、本発明ではより優れた耐衝突特性の付与を目的として９０％以上を合格とした。
【００６２】
表２の試料番号２は仕上圧延機内最高温度が本発明範囲外となっているため、マルテンサイト粒径が２μｍ超となり、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている例である。
試料番号３は仕上圧延機内最低温度が本発明範囲外となっているため、フェライト域で圧延され、粗大なフェライト結晶粒とともにマルテンサイト粒径が２μｍを超え、耐衝突特性、伸びフランジ性が劣っている例である。
試料番号４は熱間圧延後の冷却速度が本発明範囲外の２０℃／ｓと小さいため、マルテンサイト面積率は確保されるもののマルテンサイト粒径を小さく出来ず、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている例である。
試料番号５は熱延終了温度が２００℃超で本発明範囲外であるため、マルテンサイト面積率を５％以上確保出来ず、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている例である。
【００６３】
試料番号６、２５、３６はスラブの化学成分が本発明の範囲外となっている例である。試料番号６はＣが本発明の範囲外となっている例であり、０．０４％未満であるためにマルテンサイト面積率を５％以上確保出来きず、その結果、十分な降伏点上昇量やＢＨ量を得ることが出来ていない例である。
試料番号２５はＮが本発明範囲外となっている例であり、０．００５％を超えているため、析出した窒化物が鋼板組織を不均一にし、微細なマルテンサイト粒と粗大なマルテンサイト粒を共存させ、最大のマルテンサイト粒径が２μｍを超える。そのため、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている。
試料番号３６はＴｉとＮｂの合計量が本発明範囲外となっている例であり、伸びフランジ性は良好であるが、０．０５％を超えているため十分なＢＨ量が確保出来ていない例である。
【００６４】
試料番号１１、３０は焼鈍温度が本発明で規定する範囲よりも低いため、試料番号１２は焼鈍温度が高すぎるのでγ相とα相の二相共存域ではなくγ相域で焼鈍されているため、マルテンサイト粒径が大きくなり、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている。
試料番号２２、２９は冷延率が本発明範囲外となっている例であり、不均一な組織になるため、最大のマルテンサイト粒径が２μｍを超え、ＢＨ量が低く、耐衝突特性、伸びフランジ性が劣っている。
試料番号１３、３１は焼鈍後の冷却速度が本発明の範囲外となっている例で、パーライト変態が生じたため、マルテンサイト面積率が確保できず、ＢＨ量が低く、耐衝突特性が劣っている。
【００６５】
試料番号１４、２３は焼鈍後の冷却終了温度が本発明の範囲外となっている例であり、パーライト変態が生じたため、マルテンサイト面積率が確保出来ず、ＢＨ量が低く、耐衝突特性が劣っている。
試料番号１５、１６は、それぞれ溶融亜鉛めっき処理後の冷却速度、冷却終了温度が本発明の範囲外の例であり、パーライト変態が生じるため、マルテンサイト面積率が確保出来ず、ＢＨ量が低く、耐衝突特性に劣っている。
【００６６】
【表１】

【００６７】
【表２】

【００６８】
【表３】

【００６９】
前述した例以外は本発明例であり、表２からも判るように、何れの例も耐衝突特性、形状凍結性に優れた高強度鋼板となっている。
なお、試料番号２６は、（ＹＰ＋ＷＨ＋ＢＨ）／ＴＳ比が８８％であり、合格ラインの９０％に達しないものの、従来品よりも良好な耐衝突特性を示す参考例である。
【００７０】
【発明の効果】
以上説明したように、本発明によればプレス成形時の優れた伸びフランジ性とプレス成形時の優れた塗装焼付け硬化性及び耐衝突特性に優れた４４０〜７８０ＭＰａ級の高強度鋼板、高強度溶融亜鉛めっき鋼板並びに高強度合金化溶融亜鉛めっき鋼板を利用できるようになり、産業上極めて大きな意義をもつ。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a steel sheet useful for automobile body members that are mainly subjected to press baking and then subjected to paint baking, and in particular, good stretch flangeability during press molding, and paint bake hardening after press molding. The present invention relates to a high-strength steel sheet, a high-strength hot-dip galvanized steel sheet, a high-strength galvannealed steel sheet, and methods for producing them.
[0002]
[Prior art]
In recent years, from the viewpoint of global environmental conservation, the amount of carbon dioxide emission has been regulated, and improvement in fuel consumption has been desired by reducing the weight of automobiles. In order to reduce the weight of an automobile, it is an effective means to reduce the thickness of a steel plate that occupies a large weight ratio in the automobile body. Moreover, from the viewpoint of occupant protection at the time of collision, it is important to have a structure that can absorb energy at the time of collision by devising the vehicle body structure, and the development of steel materials suitable for it is important.
Under such circumstances, the usage rate of high-strength steel sheets with high yield strength at the time of collision is increasing at the same time as thinning, centering on automobile structural members and reinforcing members.
[0003]
Steel sheets for application to automobile body frame members are required to have high characteristics known as collision resistance characteristics, and specifically, a tensile strength of about 440 to 780 MPa is required. It is known that increasing the yield point (YP) and increasing the bake hardenability (BH property) are effective in improving the impact resistance of the steel sheet.
However, as the tensile strength of the steel plate increases, the formability becomes more difficult, and the press formability becomes worse, for example, the part that undergoes stretch flange forming during forming breaks. Therefore, a steel sheet having practical impact resistance characteristics without impairing stretch flangeability is desired.
[0004]
Under such circumstances, as disclosed in JP-A-6-73497, the steel sheet structure is a three-phase composite structure of ferrite, martensite, and bainite, which is a characteristic feature of the ferrite-martensite composite structure. Even if the elongation and low yield ratio (ratio of yield point to tensile strength) are ensured and the bainite structure coexists with this, even if the hard martensite phase at the origin of voids, which causes deterioration of stretch flange deformation, is reduced, Techniques that can ensure strength and also impart BH properties have been proposed.
However, the amount of BH in the technique described in this publication is 51 MPa at the maximum, and a steel sheet provided with the amount of BH is desired in order to further improve the collision resistance.
[0005]
[Problems to be solved by the invention]
Thus, so far, no high-strength steel sheet that satisfies both the BH amount necessary for satisfying sufficient collision resistance and good stretch flangeability has been known.
In view of such a current situation, the present invention aims to provide a high-strength steel sheet excellent in stretch flangeability and impact resistance.
It is also an object of the present invention to provide a high-strength hot-dip galvanized steel sheet and a high-strength galvannealed steel sheet that are excellent in stretch flangeability and impact resistance.
[0006]
[Means for Solving the Problems]
As a result of studying the yield strength and stretch flangeability of various thin steel sheets after processing, painting and baking, the inventors have obtained the following knowledge, particularly when developing high-strength steel sheets as structural members for automobiles. It was. That is, stretch flangeability correlates with the particle size of hard martensite in contact with the soft ferrite phase, and when the martensite particle size is 2 μm or less, the role as the origin of voids governing stretch flange molding is reduced. I understood. In addition, the presence of fine martensite in the steel sheet structure also improves the coating bake hardening amount (BH property), and at least 60 MPa or more can be obtained.
[0007]
More specifically, the inventors define the composition of the C, Si, Mn, P, S, Al, and N components of the steel sheet within a specific range, with the ferrite phase as the main phase, and the second phase. When the maximum particle size of martensite is 2 μm or less and the area ratio of the martensite phase is 5% or more, the desired high strength steel sheet with excellent stretch flangeability and impact resistance can be obtained. I found. The present invention is based on such new findings, and the gist thereof is as follows.
[0008]
  (1) Expressed in mass%, C is 0.04 to 0.22%, Si is 1.0% or less, Mn is 3.0% or less, P is 0.05% or less, and S is 0.01%. Hereinafter, Al is contained in an amount of 0.01 to 0.1% and N is contained in an amount of 0.001 to 0.005%, and the balance is composed of Fe and unavoidable impurities.In addition, it is composed of a ferrite phase that is the main phase and a martensite phase that is the second phase.The maximum particle size is 2μm or less,A high-strength steel sheet excellent in stretch flangeability and impact resistance, characterized in that the area ratio is 5% or more.
[0009]
(2) The composition contains one or more selected from Nb, Ti and V in a total amount of 0.008% or more and 0.05% or less, and is excellent in stretch flangeability and impact resistance as described in (1). High strength steel plate.
[0010]
(3) A high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance properties, comprising the steel sheet according to either (1) or (2) above and a hot-dip galvanized layer formed on the surface thereof.
[0011]
(4) A high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance properties, comprising the steel sheet according to either (1) or (2) above and an alloyed hot-dip galvanized layer formed on the surface thereof. .
[0012]
  (5) A slab satisfying the component composition described in either (1) or (2) above is cast and solidified and then heated at 1150 ° C. or higher for 60 minutes or longer, and finish hot rolling is performed.Temperature in finishing hot rolling millA method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability and impact resistance, characterized by performing at 870 ° C. or more and 1020 ° C. or less and cooling to 200 ° C. or less at a cooling rate of 40 ° C./s or more.
[0013]
(6) After casting and solidifying the slab satisfying the component composition according to any one of (1) and (2) above, heating at 1150 ° C. or more for 60 minutes or more, and finish hot rolling at 870 ° C. or more and 1020 ° C. or less. After the descaling process, after performing the cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature that satisfies the following formula (I) for 3 seconds or more, and then 2 to 200 ° C./s. A method for producing a high-strength cold-rolled steel sheet excellent in stretch flangeability and impact resistance, characterized by cooling at a cooling rate of less than 400 ° C.
[0014]
[Formula 6]

[0015]
(7) After casting and solidifying a slab satisfying the component composition described in either (1) or (2) above, heat at 1150 ° C. or more for 60 minutes or more, and finish hot rolling at 870 ° C. or more and 1020 ° C. or less. And after descaling, after annealing for 3 seconds or more at an annealing temperature satisfying the following formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, and then hot dip galvanizing is performed. A method for producing a high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance characteristics, characterized in that it is then cooled to 250 ° C. or lower at a cooling rate of 5 ° C./s or higher.
[0016]
[Expression 7]

[0017]
(8) A slab satisfying the component composition described in either (1) or (2) above is cast and solidified and then heated at 1150 ° C. or more for 60 minutes or more, and finish hot rolling at 870 ° C. or more and 1020 ° C. or less. And after descaling, after annealing for 3 seconds or more at an annealing temperature satisfying the following formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, and then hot dip galvanizing is performed. It is excellent in flangeability and impact resistance, characterized in that it is then held for 5 to 90 seconds in a temperature range of 450 to 550 ° C and then cooled to 250 ° C or less at a cooling rate of 5 ° C / s or more. A method for producing high strength galvannealed steel sheets.
[0018]
[Equation 8]

[0019]
(9) After casting and solidifying the slab satisfying the component composition according to any one of (1) and (2) above, heating at 1150 ° C. or more for 60 minutes or more, and finishing hot rolling at 870 ° C. or more and 1020 ° C. or less. After the descaling process, after performing the cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature that satisfies the following formula (I) for 3 seconds or more, and then 2 to 200 ° C./s. Excellent stretch flangeability and impact resistance characteristics, characterized by cooling to 350 to 500 ° C. at a cooling rate of 1, followed by hot dip galvanizing, followed by cooling to 250 ° C. or less at a cooling rate of 5 ° C./s or more Manufacturing method of high strength hot dip galvanized steel sheet.
[0020]
[Equation 9]

[0021]
(10) After casting and solidifying a slab satisfying the component composition described in either (1) or (2) above, heat at 1150 ° C. or more for 60 minutes or more, and finish hot rolling at 870 ° C. or more and 1020 ° C. or less. After the descaling process, after performing the cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature that satisfies the following formula (I) for 3 seconds or more, and then 2 to 200 ° C./s. At a cooling rate of 350 ° C. to 500 ° C., followed by hot dip galvanization, after which it is held in a temperature range of 450 to 550 ° C. for 5 to 90 seconds, and then at a cooling rate of 5 ° C./s or more and A method for producing a high-strength alloyed hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance characteristics, characterized by cooling below.
[0022]
[Expression 10]

[0023]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be described in detail.
In the high-strength steel sheet excellent in stretch flangeability and impact resistance characteristics of the present invention, the component composition is expressed by mass%, C is 0.04 to 0.22%, Si is 1.0% or less, Mn Is 3.0% or less, P is 0.05% or less, S is 0.01% or less, Al is 0.01 to 0.1%, and N is 0.001 to 0.005%, The balance is Fe and inevitable impurities.
[0024]
The C component in the steel sheet is an element important for the formation of martensite and an indispensable element for obtaining the high paint bakeability obtained with the steel of the present invention. If C is less than 0.04%, a martensite area ratio of 5% or more cannot be obtained, and the lower limit of C is 0.04%. Further, when C exceeds 0.22%, the weldability is remarkably deteriorated, so the upper limit of C is set to 0.22%.
[0025]
The Si component is an element effective for enhancing the mechanical strength of the steel sheet without impairing the ductility. When Si is added in excess of 1.0%, the plating adhesion is impaired, so this is the upper limit.
[0026]
The Mn component is an austenite stabilizing element, and is an element effective for producing a transformation product and increasing the mechanical strength of the steel sheet. However, when the Mn content exceeds 3.0%, not only melting becomes difficult, but workability deteriorates, so the upper limit of Mn is set to 3.0%.
[0027]
The P component is an element that increases the mechanical strength of the steel sheet at a low cost. However, when the P content exceeds 0.05%, brittle fracture after processing tends to occur, so the upper limit of P is set to 0.05%.
[0028]
S component exists in inclusions in steel. When the S content exceeds 0.01%, the amount of inclusions increases and the stretch flangeability deteriorates, so the upper limit of S is made 0.01%.
[0029]
The Al component is an element used for deoxidation. Therefore, if it is less than 0.01%, deoxidation becomes insufficient and workability deteriorates, so this is the lower limit. Further, if Al exceeds 0.1%, the cost is increased and the deoxidation effect is saturated. Therefore, the upper limit is made 0.1%.
[0030]
N component, like C, is an important element in order to obtain martensite formation and paint bakeability. However, when Al, which is a deoxidizing element, is present, N is present in the steel as a nitride and deteriorates ductility. Therefore, the smaller the amount of N, the better. However, to make it less than 0.001%, a de-N process is required and the manufacturing cost increases, so this is the lower limit. On the other hand, if the N content exceeds 0.005%, the deposited nitride makes the steel sheet structure non-uniform, so that a stable strength cannot be obtained, making it unsuitable for industrial production. Therefore, the upper limit of the N amount is set to 0.005%.
[0031]
The high-strength steel sheet of the present invention can further contain one or more selected from Nb, Ti and V in addition to the above-described components under a total condition of 0.008% to 0.05%.
Nb, Ti, and V are all carbonitride-forming elements, and are effective elements for forming precipitates in the steel and increasing the mechanical strength of the steel sheet, particularly the yield point. When the total amount of one or more selected from Nb, Ti, and V is less than 0.008%, the effect of increasing the yield point is hardly obtained, and the lower limit is made 0.008%. Moreover, when the total amount thereof exceeds 0.05%, workability deteriorates and precipitates in combination with C and N in steel, which is the source of paint bake hardenability, and high BH properties cannot be obtained. . Therefore, the upper limit is made 0.05%.
[0032]
The steel sheet of the present invention is composed of a ferrite phase as a main phase and a martensite phase as a second phase. If the particle size of the martensite as the second phase exceeds 2 μm, voids are likely to occur at the interface between the soft ferrite phase and the hard martensite, and good stretch flangeability cannot be obtained. Therefore, the particle size of martensite is 2 μm or less.
[0033]
At the same time, the area ratio of the martensite phase is important. If the area ratio is 5% or less, even if the martensite particle size is 2 μm or less, the introduction amount of dislocations around the martensite is small because the absolute amount of the hard martensite is small. As a result, the yield point increase amount, that is, the BH amount is also reduced. Therefore, the martensite area ratio is 5% or more.
[0034]
Here, the grain size of the martensite is the cut length of the martensite crystal when the region of 100 μm × 100 μm is observed with a scanning electron microscope at a quarter of the plate thickness of an arbitrary plate cross section. It is the maximum value. In addition, the martensite area ratio means the area ratio occupied by the martensite phase in the 100 μm × 100 μm region in the above-described structure observation.
[0035]
The high-strength steel plate of the present invention may be manufactured as a hot-rolled steel plate or a cold-rolled steel plate as long as the above conditions are satisfied. The high-strength steel plate of the present invention may be manufactured as a plated steel plate, and particularly preferably manufactured as a hot-dip galvanized steel plate or an alloyed hot-dip galvanized steel plate.
[0036]
Next, the manufacturing method of the high strength steel plate, the high strength hot dip galvanized steel plate, and the high strength alloyed hot dip galvanized steel plate in the present invention will be described in detail.
[0037]
First, in order to produce the high-strength steel sheet of the present invention as a hot-rolled steel sheet, a slab satisfying the component composition in the present invention is cast and solidified and then heated at 1150 ° C. or more for 60 minutes or more, and finish hot rolling is performed at 870 ° C. or more. It is performed at 1020 ° C. or lower, and is cooled to 200 ° C. or lower at a cooling rate of 40 ° C./s or higher.
[0038]
The slab is heated at 1150 ° C. or more for 60 minutes or more. When the heating temperature is less than 1150 ° C, or when the heating time is less than 60 minutes, a temperature distribution occurs in the slab, a rolling temperature difference occurs in the length direction and the width direction, and the dimensional accuracy and the structure in steel are reduced. Uneven and not suitable for industrial production.
[0039]
The finish hot rolling needs to pass through the finish hot rolling mill within a temperature range of 870 ° C. or more and 1020 ° C. or less. If it exceeds 1020 ° C., the austenite crystal grains become coarse, the austenite grain size after finish hot rolling also increases, and martensite of 2 μm or less cannot be obtained. Further, at 870 ° C. or lower, since a ferrite phase is generated during rolling, a uniform and fine austenite structure cannot be obtained, so that a martensite particle size locally exceeding 2 μm is generated. Further, when the cooling rate after hot rolling is less than 40 ° C./s, it is difficult to ensure the amount of martensite transformation, and even if it can be ensured, the martensite that is transformed by the growth of austenite crystal grains during cooling The particle size exceeds 2 μm. Therefore, 40 ° C./s or more is necessary. When the cooling end temperature exceeds 200 ° C., the martensite area ratio is not ensured as described above. Therefore, the cooling end temperature is set to 200 ° C. or lower.
[0040]
In order to produce the high-strength steel sheet of the present invention as a cold-rolled steel sheet, a slab satisfying the component composition in the present invention is cast and solidified and then heated at 1150 ° C. or more for 60 minutes or more, and finish hot rolling is performed at 870 ° C. or more and 1020 ° C. After the descaling process, the cold rolling process is performed at a cold rolling rate of 40% or more, and annealing is performed at an annealing temperature that satisfies the following formula (I) for 3 seconds or more, and then 2 to 200 ° C./s. At a cooling rate of less than 400 ° C.
[0041]
## EQU11 ##

[0042]
If the cold rolling rate is less than 40%, cold strain is not uniformly introduced to the center of the plate thickness, and the amount of cold strain in the center layer after annealing is small. Therefore, abnormal ferrite grains in which adjacent cold-rolled crystal grains grow together grow from the ferrite grains produced by recrystallization, so that the austenite grain size produced by recrystallization also becomes non-uniform. Therefore, the grain size of martensite that is transformed from austenite grains is not sufficiently small, and sufficient stretch flangeability cannot be obtained.
[0043]
In the annealing process after cold rolling, it is important to heat and hold in a two-phase region of α phase and γ phase that satisfies the formula (I). In the temperature range above the upper limit of the formula (I), since it becomes a γ phase single phase, the nucleation frequency of the ferrite phase generated by the γ / α transformation decreases, so the ferrite grain size becomes coarse and at the same time austenite grains grow. Therefore, a martensite particle size of 2 μm or less cannot be obtained. Further, in the temperature range below the lower limit of the formula (I), the austenite grains grow at the same time as the ferrite crystal grains grow, but the martensite is desired. The particle size cannot be obtained.
[0044]
Further, when the holding time is less than 3 seconds, the structure is likely to be non-uniform, which leads to variations in material in the longitudinal direction or the width direction of the steel sheet. Therefore, the holding time is set to 3 seconds or more. When the cooling rate after annealing in the two-phase region is less than 2 ° C./s, pearlite transformation occurs, and it becomes difficult to secure the amount of martensite, and when it exceeds 200 ° C./s, the steel plate shape is significantly deteriorated. 2 to 200 ° C./s. The cooling end temperature is preferably less than 400 ° C. This is because when the temperature is 400 ° C. or higher, a pearlite transformation phase is generated, and a sufficient martensite phase area ratio cannot be obtained.
[0045]
In order to produce the high-strength steel sheet of the present invention as a hot-dip galvanized steel sheet, as described above, the high-strength steel sheet of the present invention is (1) a hot-rolled steel sheet and then hot-dip galvanized. (2 ) After forming a cold-rolled steel sheet, any method of hot dip galvanizing treatment can be employed.
[0046]
In the case of (1), a slab satisfying the component composition in the present invention is cast and solidified, heated at 1150 ° C. or more for 60 minutes or more, finish hot rolling is performed at 870 ° C. or more and 1020 ° C. or less, and then descaling treatment is performed. Then, after annealing for 3 seconds or more at an annealing temperature satisfying the above formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, then hot dip galvanized, and then 5 Cool to 250 ° C. or less at a cooling rate of at least ° C./s.
[0047]
In the case of (2) above, the slab satisfying the component composition in the present invention is cast and solidified, then heated at 1150 ° C. or more for 60 minutes or more, finish hot rolling is performed at 870 ° C. or more and 1020 ° C. or less, and after descaling treatment After performing the cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature that satisfies the above formula (I) for 3 seconds or more, and then at a cooling rate of 2 to 200 ° C./s. It is cooled to 500 ° C., then hot dip galvanized, and then cooled to 250 ° C. or lower at a cooling rate of 5 ° C./s or higher.
[0048]
In the annealing process after descaling or subsequent cold rolling,
The reason for heating and holding at an annealing temperature satisfying the formula (I) is as described above.
[0049]
The cooling rate after annealing in the two-phase region is 2 to 200 ° C./s. If it is less than 2 ° C./s, pearlite transformation occurs, making it difficult to ensure the amount of martensite, and if it exceeds 200 ° C./s, the shape of the steel sheet is significantly deteriorated.
[0050]
The cooling end temperature is 350 ° C to 500 ° C. If the cooling end temperature is less than 350 ° C., the zinc bath temperature is lowered, which increases the cost such as heating the bath temperature. If it exceeds 500 ° C., the bainite transformation occurs from the γ phase, and the desired martensite area ratio is I can't get it.
[0051]
In the cooling step after plating, when the cooling rate does not reach less than 5 ° C./s or when cooling is not performed to 250 ° C. or lower, a pearlite transformation phase is generated and a martensite phase cannot be obtained. Therefore, it is cooled to 250 ° C. or lower at 5 ° C./s or higher.
Here, the hot dip galvanizing method is not particularly limited, and for example, a commonly used method such as using an in-line annealing method of continuous hot dip galvanizing equipment may be used.
[0052]
In order to produce the high-strength steel sheet of the present invention as an alloyed hot-dip galvanized steel sheet, (1) hot-rolled steel sheet is formed, and then hot-dip galvanized treatment is performed, Any of the following methods can be employed: a method of alloying, (2) a method of producing a cold-rolled steel sheet, performing a hot dip galvanizing process, and further performing an alloying process.
[0053]
In the case of (1), a slab satisfying the component composition in the present invention is cast and solidified, heated at 1150 ° C. or more for 60 minutes or more, finish hot rolling is performed at 870 ° C. or more and 1020 ° C. or less, and then descaling treatment is performed. Then, after annealing for 3 seconds or more at an annealing temperature satisfying the above formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, then hot dip galvanized, and then 450 After holding for 5 to 90 seconds in a temperature range of ˜550 ° C., cooling to 250 ° C. or less at a cooling rate of 5 ° C./s or more.
[0054]
In the case of (2) above, the slab satisfying the component composition in the present invention is cast and solidified, then heated at 1150 ° C. or more for 60 minutes or more, finish hot rolling is performed at 870 ° C. or more and 1020 ° C. or less, and after descaling treatment After performing the cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature that satisfies the above formula (I) for 3 seconds or more, and then at a cooling rate of 2 to 200 ° C./s. After cooling to 500 ° C., hot dip galvanizing is performed, and thereafter, holding in a temperature range of 450 to 550 ° C. for 5 to 90 seconds, cooling is performed to 250 ° C. or less at a cooling rate of 5 ° C./s or more.
[0055]
The production of the alloyed hot-dip galvanized steel sheet is the same as that of the hot-dip galvanized steel sheet except that the alloying treatment is performed subsequent to the hot-dip galvanizing treatment.
[0056]
In the alloying treatment, it is necessary to hold at a heating temperature of 450 to 550 ° C. for 5 to 90 seconds.
When the holding in the temperature range of 450 to 550 ° C. is less than 5 seconds, a sufficient alloy layer of iron and zinc cannot be formed, and characteristics as an alloyed hot-dip galvanized steel sheet cannot be obtained. Also, holding for more than 90 seconds in the above temperature range generates iron carbide cementite (Fe3C) at the ferrite grain boundary, so the amount of C interacting with dislocations introduced by press working is reduced, and the amount of BH Decrease. Therefore, the holding time is in the range of 5 to 90 seconds.
[0057]
The method of applying the alloyed hot dip galvanizing is not particularly limited, and for example, a conventional method such as using an in-line annealing method of continuous hot dip galvanizing equipment may be used.
[0058]
【Example】
Next, the present invention will be further described with reference to examples. Needless to say, the present invention is not limited to these examples.
[0059]
A steel having the component composition shown in Table 1 was cast and solidified, or after being cooled to room temperature, and then hot-rolled under the conditions shown in Table 2 to produce a hot-rolled steel sheet. The steel slab was heated at 1150 ° C. or more for 60 minutes or more and then hot-rolled. A part of these hot-rolled steel sheets was further descaled, and then subjected to hot dip galvanizing or galvannealed galvanizing under the conditions shown in Table 2. Alternatively, another part of the hot-rolled steel sheet was descaled and then cold-rolled under the conditions shown in Table 2 and annealed to produce a cold-rolled steel sheet. A part of the cold-rolled steel sheet thus obtained was subjected to an annealing treatment followed by a hot dip galvanizing treatment or an alloyed hot dip galvanizing treatment under the conditions shown in Table 2.
[0060]
While microscopically observing the microstructure of each steel plate obtained, specimens were taken perpendicular to the rolling direction, and yield strength (YP), tensile strength (TS), paint bake hardening (BH) amount by tensile test I investigated. Here, the yield strength and the tensile strength were measured according to the method described in JIS Z 2241 using No. 5 test piece described in JIS Z 2201. Further, the BH amount is an increase amount from a stress after 2% tensile pre-deformation to a yield stress when re-tensioned after aging treatment at 170 ° C. for 20 minutes. In addition, a 200 mm square test piece was cut out, a hole expanding test using a flat bottom cylindrical punch (diameter 100 mm) was performed on a hole having a hole diameter of 25 mm, and stretch flangeability was evaluated. Stretch flangeability is evaluated by the hole expansion ratio λ defined by (hole diameter when cracking occurs at the hole edge−initial hole diameter) ÷ initial hole diameter. did.
[0061]
  The above results are shown in Table 3. In Table 3, YP + WH + BH is the amount obtained by adding the strength increase (WH) from the yield strength due to the pre-strain deformation of 2% and the paint bake hardening amount (BH) to the yield strength (YP). After yield stress("Yield stress when re-tensioned" described in the measurement of BH amount above)It corresponds to. Therefore, the larger this value, the greater the yield strength at the time of collision, and the better the collision resistance. In the present invention, the ratio of (YP + WH + BH) to TS, that is, (YP + WH + BH) / TS, is defined as an index representing the impact resistance characteristics after pressing and paint baking of members having the same formability on the original plate. In the conventional steel plate, it is about 85% at the highest, and in the present invention, 90% or more was accepted for the purpose of imparting more excellent collision resistance.
[0062]
Sample No. 2 in Table 2 is an example in which the maximum temperature in the finishing mill is outside the range of the present invention, so the martensite particle size exceeds 2 μm, the BH amount is low, and the impact resistance and stretch flangeability are inferior. is there.
Sample No. 3 is rolled in the ferrite region because the minimum temperature in the finish rolling mill is outside the range of the present invention, and the martensite grain size exceeds 2 μm together with coarse ferrite crystal grains, resulting in poor impact resistance and stretch flangeability. This is an example.
Sample No. 4 has a cooling rate after hot rolling as small as 20 ° C./s outside the range of the present invention, so although the martensite area ratio is ensured, the martensite particle size cannot be reduced, the BH amount is low, and the collision resistance This is an example of poor properties and stretch flangeability.
Sample No. 5 has a hot rolling end temperature of over 200 ° C. and is outside the scope of the present invention, so it cannot secure a martensite area ratio of 5% or more, has a low BH content, and has poor impact resistance and stretch flangeability. It is.
[0063]
Sample numbers 6, 25 and 36 are examples in which the chemical composition of the slab is outside the scope of the present invention. Sample No. 6 is an example in which C is outside the scope of the present invention, and since it is less than 0.04%, the martensite area ratio cannot be ensured to be 5% or more. This is an example in which the BH amount cannot be obtained.
Sample No. 25 is an example in which N is outside the scope of the present invention and exceeds 0.005%, so that the deposited nitride makes the steel sheet structure non-uniform, and fine martensite grains and coarse martensite. Coexists with grains, and the maximum martensite particle size exceeds 2 μm. Therefore, the amount of BH is low, and the impact resistance and stretch flangeability are inferior.
Sample No. 36 is an example in which the total amount of Ti and Nb is outside the scope of the present invention, and the stretch flangeability is good, but since it exceeds 0.05%, a sufficient amount of BH cannot be secured. It is an example.
[0064]
Since Sample Nos. 11 and 30 have an annealing temperature lower than the range specified in the present invention, Sample No. 12 is annealed in the γ phase region, not in the two-phase coexistence region of the γ phase and the α phase because the annealing temperature is too high. For this reason, the martensite particle size becomes large, the amount of BH is low, and the impact resistance and stretch flangeability are inferior.
Sample Nos. 22 and 29 are examples in which the cold rolling rate is outside the scope of the present invention, and because the structure becomes non-uniform, the maximum martensite particle size exceeds 2 μm, the BH amount is low, Stretch flangeability is inferior.
Sample Nos. 13 and 31 are examples in which the cooling rate after annealing is outside the scope of the present invention, because pearlite transformation occurred, the martensite area ratio could not be secured, the BH amount was low, and the impact resistance characteristics were inferior. Yes.
[0065]
Sample Nos. 14 and 23 are examples in which the cooling end temperature after annealing is outside the range of the present invention, and since pearlite transformation occurred, the martensite area ratio could not be ensured, the BH amount was low, and the impact resistance characteristics were Inferior.
Sample numbers 15 and 16 are examples in which the cooling rate after the hot dip galvanizing treatment and the cooling end temperature are outside the scope of the present invention, and pearlite transformation occurs. Therefore, the martensite area ratio cannot be secured and the BH amount is low. Inferior to collision resistance.
[0066]
[Table 1]

[0067]
[Table 2]

[0068]
[Table 3]

[0069]
  Examples other than those described above are examples of the present invention, and as can be seen from Table 2, each example is a high-strength steel sheet having excellent impact resistance and shape freezing properties.
  Sample No. 26 is a reference example showing a better collision resistance than the conventional product, although the (YP + WH + BH) / TS ratio is 88% and does not reach 90% of the acceptable line.
[0070]
【The invention's effect】
As described above, according to the present invention, a high-strength steel sheet of 440 to 780 MPa class excellent in stretch flangeability during press molding, excellent paint bake hardenability and impact resistance during press molding, high-strength melting Zinc-coated steel sheets and high-strength alloyed hot-dip galvanized steel sheets can be used, which has a great industrial significance.

Claims (10)

Expressed in mass%, C is 0.04 to 0.22%, Si is 1.0% or less, Mn is 3.0% or less, P is 0.05% or less, S is 0.01% or less, Al the 0.01 to 0.1% and N containing from 0.001 to 0.005 percent, while have a component composition and the balance Fe and unavoidable impurities, and a ferrite phase as a main phase, a second phase A high-strength steel sheet excellent in stretch flangeability and impact resistance, characterized in that it is composed of a certain martensite phase and has a maximum grain size of 2 μm or less and an area ratio of 5% or more. The high-strength steel sheet excellent in stretch flangeability and impact resistance properties according to claim 1, wherein the component composition contains at least one selected from Nb, Ti and V in a total amount of 0.008% to 0.05%. A high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance properties, comprising the steel plate according to claim 1 and a hot-dip galvanized layer formed on the surface thereof. A high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance properties, comprising the steel plate according to claim 1 and an alloyed hot-dip galvanized layer formed on the surface thereof. A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at a temperature in a finish hot rolling mill of 870 ° C or higher and 1020. A method for producing a high-strength hot-rolled steel sheet excellent in stretch flangeability and impact resistance characteristics, characterized by being carried out at a temperature of ℃ or less and cooled to a temperature of 200 ℃ or less at a cooling rate of 40 ℃ / s. A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at 870 ° C or higher and 1020 ° C or lower, and descaling is performed. After the treatment, after performing a cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature satisfying the following formula (I) for 3 seconds or more, and then at a cooling rate of 2 to 200 ° C./s. A method for producing a high-strength cold-rolled steel sheet excellent in stretch flangeability and impact resistance, characterized by cooling to less than 400 ° C.

A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at 870 ° C or higher and 1020 ° C or lower, and descaling is performed. After the treatment, after annealing for 3 seconds or more at an annealing temperature satisfying the following formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, followed by hot dip galvanization, A method for producing a high-strength hot-dip galvanized steel sheet excellent in stretch flangeability and impact resistance, characterized by cooling to 250 ° C or lower at a cooling rate of 5 ° C / s or higher.

A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at 870 ° C or higher and 1020 ° C or lower, and descaling is performed. After the treatment, after annealing for 3 seconds or more at an annealing temperature satisfying the following formula (I), it is cooled to 350 to 500 ° C. at a cooling rate of 2 to 200 ° C./s, followed by hot dip galvanization, A high-strength alloy excellent in flangeability and impact resistance, characterized by holding for 5 to 90 seconds in a temperature range of 450 to 550 ° C. and then cooling to 250 ° C. or less at a cooling rate of 5 ° C./s or more. Method for producing a galvannealed steel sheet.

A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at 870 ° C or higher and 1020 ° C or lower, and descaling is performed. After the treatment, after performing a cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature satisfying the following formula (I) for 3 seconds or more, and then at a cooling rate of 2 to 200 ° C./s. Cooling to 350-500 ° C, followed by hot dip galvanization, followed by cooling to 250 ° C or lower at a cooling rate of 5 ° C / s or higher, high strength melting with excellent stretch flangeability and impact resistance Manufacturing method of galvanized steel sheet.

A slab satisfying the component composition according to claim 1 or 2 is cast and solidified and then heated at 1150 ° C or higher for 60 minutes or longer, and finish hot rolling is performed at 870 ° C or higher and 1020 ° C or lower, and descaling is performed. After the treatment, after performing a cold rolling process at a cold rolling rate of 40% or more, annealing is performed at an annealing temperature satisfying the following formula (I) for 3 seconds or more, and then at a cooling rate of 2 to 200 ° C./s. Cool to 350 to 500 ° C., then apply hot dip galvanizing, and then hold for 5 to 90 seconds in a temperature range of 450 to 550 ° C., then cool to 250 ° C. or less at a cooling rate of 5 ° C./s or more. A method for producing a high-strength alloyed hot-dip galvanized steel sheet having excellent stretch flangeability and impact resistance characteristics.