JP3924159B2 - High-strength thin steel sheet with excellent delayed fracture resistance after forming, its manufacturing method, and automotive strength parts made from high-strength thin steel sheet - Google Patents
High-strength thin steel sheet with excellent delayed fracture resistance after forming, its manufacturing method, and automotive strength parts made from high-strength thin steel sheet Download PDFInfo
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Description
【０００１】
【発明の属する技術分野】
本発明は、特に高強度薄鋼板において問題となる、置き割れや遅れ破壊を抑制した高強度薄鋼板及びその製造方法並びにそれらにより作成される自動車用強度部品に関するものである。
【０００２】
【従来の技術】
従来、ボルト、ＰＣ鋼線やラインパイプといった用途には高強度鋼が多く使われており、７８０ＭＰａ以上の強度になると、鋼中への水素の侵入により遅れ破壊が発生することが知られている。これに対し、▲１▼薄鋼板は板厚が薄いため水素が侵入しても短時間で放出されること、▲２▼加工性の点で７８０ＭPa以上の鋼板の利用がほとんどなかったことなどから、遅れ破壊に対する問題意識は低かったと言える。
【０００３】
しかし、最近では自動車の軽量化や衝突安全性の向上の必要性から、バンパーやインパクトビーム等の補強材やシートレール等に７８０ＭＰａ以上の超高強度鋼板をプレス成形やパイプ成形や曲げ加工や端面加工や穴拡げ加工などを施して、使用に供する場合が急速に増えてきている。したがって、耐遅れ破壊性を備えた超高強度薄鋼板の開発が急務である。
【０００４】
これまで、耐遅れ破壊を向上させる技術はほとんどがボルトや条鋼、厚板といった製品のままでかつ耐力または降伏応力以下で使用されることの多い鋼材に対して開発されてきた。
【０００５】
例えば条鋼・ボルト用鋼においては、焼き戻しマルテンサイトを中心に開発が行われ、「遅れ破壊解明の新展開」（（社）日本鉄鋼協会 材料の組織と特性部会 高強度鋼の遅れ破壊研究会 平成９年１月）にＣｒ、ＭｏやＶといった焼き戻し軟化抵抗性を示す添加元素が耐遅れ破壊性向上に有効であることが報告されている。これは、合金炭化物を析出させて、これを水素のトラップサイトに活用することで遅れ破壊形態を粒界から粒内破壊へと移行させる技術である。しかし、これらの鋼はＣ量０．４％以上で合金元素も多く含むことから、薄鋼板で要求される加工性や溶接性が劣悪で、さらに、合金炭化物析出には数時間以上という析出熱処理が必要なため、製造性にも問題がある。また、特開平１１―２９３３８３号でＴｉ、Ｍｇを主体とする酸化物が水素性欠陥を防ぐことに効果があるとされている。しかし、これは対象が厚鋼板であり、特に大入熱の溶接後の遅れ破壊については考慮されているものの、薄鋼板に要求される加工度の高い成形加工を受けたり、端面加工に伴うバリ発生等の遅れ破壊現象に及ぼす影響については一切考慮されていない。さらには、薄鋼板の基本的特性である加工性についての考慮も一切無い。
【０００６】
一方、薄鋼板の遅れ破壊に関しては、（例えば、山崎ら CAMP-ISIJ vol.5 p1839 〜1842 (1992))残留オーステナイト量の加工誘起変態に起因した遅れ破壊の助長について報告されている。これは、薄鋼板の成型加工を考慮したものであるが、耐遅れ破壊性を劣化させない残留オーステナイト量の規制について述べられている。すなわち、特定の組織を持つ高強度薄鋼板に関するものであり、根本的な耐遅れ破壊向上対策とは言えない。
【０００７】
【発明が解決しようとする課題】
上記のように、特に薄鋼板の使用環境や現状設備による製造性を考慮し、基本的特性である成形加工性を確保しつつ使用前の成形加工等の遅れ破壊に対する対策を講じた開発事例はほとんどない。
【０００８】
【課題を解決するための手段】
本発明者らは、以上のような背景から、薄鋼板における使用環境および現状設備での製造方法を十分に考慮して、根本的に耐遅れ破壊性を向上させる方法を見出すに至った。すなわち、Ｍｇの化合物または複合化合物形成させて、これの形態を制御することで、高強度薄鋼板の成形加工性を劣化させることなく、成形加工後の耐遅れ破壊性を向上させることを見出した。かつ現状の製造設備（熱間圧延、連続焼鈍、溶融亜鉛めっき、電気めっき設備など）を用いた前記の高強度薄鋼板の有効な製造方法を見出した。詳細は以下の通りである。
【０００９】
（１）重量％にて、
Ｃ ：０．０５％〜０．３％、
Ｓｉ：３．０％以下、
Ｍｎ：０．０１〜３．０％、
Ｐ ：０．０２％以下、
Ｓ ：０．０２％以下、
Ａｌ：０．０１％〜３．０％、
Ｎ ：０．０１％以下、
Ｍｇ：０．０００２％〜０．０１％
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とし、鋼板の組織中に残留オーステナイトが体積率で７％以下であり、Ｍｇの酸化物、Ｍｇの硫化物、Ｍｇの複合晶出物およびＭｇの複合析出物のいずれか１種以上が、
平均粒子径ｄ：０．０１〜５．０μｍ
密度ρ：１平方ｍｍあたり１００〜１０００００個
分布：平均粒子径からの標準偏差σと平均粒子径ｄの比σ / ｄ≦１．０
を満たす分布形態を有し、さらに、前記平均粒子径ｄ、前記密度ρ、残留オーステナイトの体積率Ｖγ ( ％ ) 、引張強度ＴＳ ( ＭＰａ ) が式（Ａ）を満たすことを特徴とする、成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
式（Ａ）：１０００（Ｖγ−０．１）−５．５＋α（Ｍｇ−４０)２−５０（ｄ−０．２)２＋１．１ｌｎρ＋７００（ＴＳ−６８０）−０．９≧１０
＊）α＝−０．００５（Ｍｇ≦４０）、α＝−０．００２（Ｍｇ＞４０）
Ｖγ：残留オーステナイト体積率（％）
Ｍｇ：Ｍｇ量（重量ｐｐｍ）
ｄ ：平均粒子径（μｍ）
ρ ：密度（個／ｍｍ２）
ＴＳ：引張強度（ＭＰａ）
式（Ａ）の条件：[1]１０００（Ｖγ−０．１）−５．５≧１０の時は１０００（Ｖγ−０．１）−５．５＝１０
[2]２≦Ｍｇ≦１００ｐｐｍ
[3]０．０１≦ｄ≦５．０μｍ、（ｄ−０．２）２≦０．２の時は（ｄ−０．２）２ ＝０．２
[4]１００≦ρ≦１０００００個／ｍｍ２
[5]７８０ＭＰａ≦ＴＳ
【００１０】
（２）更に、重量％にて、
Ｖ ：０．００５〜１％、
Ｔｉ：０．００２〜１％、
Ｎｂ：０．００２〜１％、
Ｚｒ: ０．００２〜１％、
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１１】
（３）更に、重量％にて、
Ｃｒ：０．００５〜５％、
Ｍｏ：０．００５〜５％、
Ｗ ：０．００５〜５％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）または（２）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１２】
（４）更に、重量％にて、
Ｃｕ：０．００５〜２．０％
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）〜（３）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１３】
（５）更に、重量％にて、
Ｎｉ：０．００５〜２．０％、
Ｃｏ：０．００５〜２．０％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）から（４）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１４】
（６）更に、重量％にて、
Ｂ ：０．０００２〜０．１％、
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）から（５）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１５】
（７）更に、重量％にて、
ＲＥＭ：０．０００５〜０．０１％、
Ｃａ：０．０００５〜０．０１％、
Ｙ ：０．０００５〜０．０１％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする（１）から（６）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１６】
（８）（１）から（７）のいずれかに記載の組成からなる鋳片を製造し、Ａｒ3 点以上の仕上温度で熱間圧延を施し、５００〜８００℃で捲取り、次いで通常の酸洗の後、圧下率を３０〜８０％として冷間圧延後、焼鈍工程で６００℃以上９５０℃以下に均熱して再結晶焼鈍を施し、次いで調質圧延を施した成形加工後の耐遅れ破壊性に優れた高強度薄鋼板の製造法。
【００１７】
（９）（８）記載の製造方法で、焼鈍後２００〜７００℃の温度域で１分から１０時間保持することを特徴とする、成形加工後の耐遅れ破壊性に優れた高強度薄鋼板の製造法。
【００１８】
（１０）（１）から（７）に記載の高強度鋼板が熱延鋼板または冷延鋼板であることを特徴とする、成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００１９】
（１１）鋼板に亜鉛めっきの表面処理を施したことを特徴とする（１）から（７）に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００２０】
（１２）（９）または（１０）に記した鋼板に、さらにフィルムラミネート処理をする成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
【００２１】
（１３）（１）から（７）に記載の高強度薄鋼板により作成された自動車用強度部品。
【００２２】
【発明の実施の形態】
焼き戻しマルテンサイト鋼では遅れ破壊は旧オーステナイト粒界等に水素が集積することにより、ボイド等が発生し、その部分が起点となって破壊を生じると考えられている。そこで、水素のトラップサイトを均等かつ微細に分散させて、その部分に水素をトラップさせると、拡散性水素濃度が下がり、遅れ破壊の感受性が下がる。前出の特開平１１―２９３３８３号にあるように、ＭｇおよびＴｉを複合添加した厚鋼板における酸化物の分散形態制御で、水素起因の耐遅れ破壊性が向上することが分かっている。しかし、薄鋼板のように成形加工を受けるため、高い残留応力が発生したり、加工端面にバリ等があると必然的に耐遅れ破壊性も劣化するため、これに伴う遅れ破壊特性の劣化を補足できない。このように、薄鋼板の使用形態を考慮した遅れ破壊特性に関する研究は少なく、ＭｇやＴｉの酸化物形態制御のみでは解決できない。また、トラップサイトの微細分散は、薄鋼板の基本特性である延性を劣化させる懸念もある。
【００２３】
本発明者らは、上述の背景を踏まえて、薄鋼板の使用環境、すなわち成形加工後においても耐遅れ破壊性を確保・向上させるため、種々の晶出物、析出物に加えて、鋼板の強度、組織の影響をそれぞれ検討した。その結果、薄鋼板の使用環境下で、高い残留応力下や端面のバリ発生があっても、耐遅れ破壊性を向上・確保するための技術を見出した。すなわち、
▲１▼ Ｍｇを含む酸化物または硫化物と、それらと複合晶出または析出した化合物の分散形態制御。
▲２▼ 鋼板のミクロ組織中の残留オーステナイト量。
▲３▼ 鋼板の強度。
をそれぞれ制御することで、有効に水素のトラップサイトであるＭｇの化合物または複合晶出・析出物を効果的に分散させ、延性および成形加工後の耐遅れ破壊性を両立させる事ができる。これを満たすための条件として、式（Ａ）を規定した。尚、▲１▼のＭｇを含む酸化物または硫化物と、それらと複合晶出または析出した化合物は、粒内（旧オーステナイト粒界等のミクロ組織の相界面を除く）にあると、より遅れ破壊特性向上に有効である。
【００２４】
上記の本発明の▲１▼▲２▼▲３▼を得るには製造条件を特定して、種々の元素の酸化物、窒化物、硫化物等の晶出物や析出物が水素のトラップサイトになり得る形態を制御することも有効である。
【００２５】
式（Ａ）を満たすことで、高強度薄鋼板の耐遅れ破壊性を成形加工後においても十分に確保できる。
これは薄鋼板の成形加工により導入される転位や残留応力場とトラップサイトとなる粒子の相互作用が、厚鋼板での熱間圧延や溶接後冷却時に導入される転位や残留応力とのそれとは異なることや、薄鋼板と厚鋼板の熱処理方法の違いに起因すると考えられる。
上記▲１▼と▲２▼は以下のように限定する。
【００２６】
残留オーステナイト量：残留オーステナイトは加工誘起変態によりマルテンサイトになると遅れ破壊感受性を大きくしてしまうため、上限を体積率で７％以下とした。
【００２７】
平均粒子径：平均粒子径は、０．０１μｍから５．０μｍに限定した。水素のトラップサイトとしてはある程度の大きさが必要であり、かつ、微細な粒子が多量にあることは薄鋼板の延性を確保する上でも好ましいものではなく、製造も困難となる。したがって、下限を０．０１μｍとした。また、粗大粒子はトラップサイトとしての作用がなくなる上破壊の起点となり得るので５．０μｍを上限とした。
【００２８】
密度：粒子密度は、１００〜１０００００個／ｍｍ2 とした。粒子密度が低いことは、トラップサイト数が少ないことを意味し、加工後の耐遅れ破壊性を確保できないため、下限を１００個／ｍｍ2 とした。また、高密での場合には、延性や成形加工性が劣化することおよび耐遅れ破壊性向上効果も飽和することから、１０００００個／ｍｍ2 を上限とした。
【００２９】
分布：粒子の分布を、平均粒子径からの標準偏差σと平均粒子径ｄの比σ / ｄが、σ / ｄ≦１．０を満たすこととした。σ／ｄ＞１．０とは、粒子分布が広範囲にわたることを意味し、耐遅れ破壊向上効果が同じ平均粒径に比べて小さくなり、延性劣化や破壊の起点数の増加にもつながることから、上限を１．０以下とした。
【００３０】
ここで、Ｍｇ化合物を含む粒子の測定について述べる。粒子の測定は、薄膜または抽出レプリカのサンプルを用いて、走査型または透過型電子顕微鏡にて、５０００〜１０００００倍の倍率で観察を行い、最低３０視野を測定することで得られる値とする。粒子径は、画像解析による円相当径にて評価する。また、密度を求める際には、複合析出または晶出物は１ヶとして数える。組成分析は、ＥＤＸおよびＥＥＬＳを用い、構造解析はDiffraction pattern を解析することで行った。各複合化合物は、Ｍｇの他合金添加元素（例えばＴｉ、Ｎｂ、Ｖ、Ｃｒ、Ｍｏ、ＲＥＭ、Ｃａなど）を含有した化合物（炭化物、窒化物、酸化物や硫化物など) である。
【００３１】
以下に本発明を更に詳細に説明する。
本発明では高強度薄鋼板について述べているが、主に引張強度で７８０ＭＰａ以上、板厚は０．５ｍｍ〜４．０ｍｍの鋼板についてである。
次に式（Ａ）であるが、残留オーステナイトの体積率、平均粒子径、密度、Ｍｇ量、引張強度が耐遅れ破壊性の要因としてあげられることに加えて、図１より設定した。式（Ａ）の左辺を関数ｆ（Ｖγ、Ｍｇ、ｄ、ρ、ＴＳ）とおくと、ｆ（Ｖγ、Ｍｇ、ｄ、ρ、ＴＳ）の値が１０以上で著しく耐遅れ破壊性が向上する。さらに各変数についての耐遅れ破壊性に対する影響を図２〜４に示す。図中の○は耐遅れ破壊性が良好であり、×は不良であったことを示す。
【００３２】
図２はf(Ｖγ）と残留オーステナイト体積率Ｖγの相関グラフである。グラフの条件はＭｇ：３０ｐｐｍ、平均粒子径：０．４μｍ、密度：１５００個／ｍｍ2 、引張強度：１４８０ＭＰａである。Ｖγが高いと耐遅れ破壊性が劣化するが、f(Ｖγ）値が高い発明鋼はＶγが７％以下で良好な耐遅れ破壊性を示す。また、Ｖγが７％以下でも×の比較鋼はＭｇ、粒径、密度が範囲外のためf(Ｖγ）＜１０となり耐遅れ破壊性が悪化したためである。
【００３３】
図３はf(Mg) とＭｇ添加量の相関グラフである。グラフの条件は残留オーステナイトの体積率：３．０％、平均粒子径：０．４μｍ、密度：１５００個／ｍｍ2 、引張強度：１４８０ＭＰａである。Ｍｇ添加量が２０〜７０ｐｐｍに特に耐遅れ破壊性が良好な領域がある。また、Ｍｇが１００ｐｐｍ以下で×のものは残留オーステナイト、粒径、密度が範囲外のためf(Mg) ＜１０となり耐遅れ破壊性が悪化したためである。
【００３４】
図４はf(ρ) と晶出物及び析出物の密度の相関グラフである。グラフの条件は残留オーステナイトの体積率：３．０％、Ｍｇ：３０ｐｐｍ、引張強度：１３８０ＭＰａである。密度が低いと耐遅れ破壊性が悪いと言える。また、密度ρが請求項の範囲内であるが×のものは、残留オーステナイト、Ｍｇ、粒径が範囲外のためf(ρ) ＜１０となり耐遅れ破壊性が悪化したためである。
以上により、式（Ａ）を満たせば耐遅れ破壊性に優れるとした。
【００３５】
次に、本発明における鋼の化学成分の限定理由について説明する。
Ｃは、鋼板の強度を上昇できる元素である。特にマルテンサイトやオーステナイトなどの硬質相を生成し高強度化には必須の元素であり、７８０ＭＰａ以上の強度を得るためには０．０５％以上が必要であるが、逆に多く含有すると、脆性破壊の起点となるセメンタイトを増加させるため、水素脆性を生じ易くする。従って、上限を０．３％とした。
【００３６】
Ｓｉは、材質を大きく硬質化する置換型固溶体強化元素であり、鋼板の強度を上昇させることに有効なうえ、セメンタイト析出を抑制する元素であるが、３．０％を超えると熱間圧延でのスケール除去にコストがかかり経済的に不利なため、３．０％を上限とする。また、添加量が多いとめっき性を劣化させるため、めっき性を向上させるためには０．６％以下が望ましい。
【００３７】
Ｍｎは、鋼板の強度上昇に有効な元素である。しかし、０．０１％未満ではこの効果が得られないので、下限値を０．０１％とした。逆に多いとＰ、Ｓとの共偏析を助長するだけでなく、加工性が劣化する場合があるため３．０％を上限値とする。
【００３８】
Ｐは、粒界偏析による粒界破壊の助長をする元素であり、低い方が望ましいが、極低下は製造コスト上好ましくない。また耐食性を劣化させる元素であるため、上限を０．０２％とする。
【００３９】
Ｓは、腐食環境下での水素吸収を助長する元素であり、低い方が望ましいが、極低下は製造コスト上好ましくない。特に加工性を高めるためには低い方が望ましく上限を０．０２％とする。
【００４０】
Ａｌは、脱酸のため０．０１％以上を添加するが、添加量が増加するとアルミナ等の介在物が増加し、加工性が劣化する及び溶接性を劣化するため３．０％を上限とする。尚、０．２％以上添加すると残留オーステナイトの生成を抑制する効果があるため望ましい。
【００４１】
Ｎは、加工性劣化や溶接時のブローホール発生にも寄与するため少ない方が良い。０．０１％を越えると加工性が劣化してくるので、０．０１％を上限とする。
【００４２】
Ｍｇは、自身の化合物が耐遅れ破壊性向上に効果的なだけでなく、他元素との複合析出物または複合昇晶出物を生成させ、かつそれらの形態を耐遅れ破壊性向上に寄与するよう制御するために必要な元素であることから、０．０００２％以上とした。しかし、０．０１％超では粗大酸化物および硫化物を生成して、形態制御に効果的でなくなる上、薄鋼板の基本的要求特性である成形加工性を低下させるため、上限を０．０１％とした。
【００４３】
次に、Ｖ、Ｔｉ、Ｎｂ、Ｚｒは強炭化物生成元素であり、析出物や介在物を生成させて強度及び耐遅れ破壊性を改善するために必要な元素である。
更にＶは、鋼板の強度上昇及び粒径の微細化に有効な元素である。しかし、０．００５％未満ではこの効果が得られないために、下限値を０．００５％とした。逆に、１％超含有すると炭窒化物の析出が顕著になり、延性低下が著しくなる。このため上限値を１％とした。
【００４４】
更にＴｉは、鋼板の強度上昇及び粒径の微細化に有効な元素である。しかし、０．００２％未満では析出物の個数が低下するために、下限値を０．００２％とした。逆に、１％超では粗大析出または昇出物が生成するために加工性および耐遅れ破壊性が低下する。このため、上限値を１％とした。
【００４５】
更にＮｂは、鋼板の強度上昇及び細粒化に有効な元素である。しかし、０．００２％未満ではこれらの効果が得られないため、下限値を０．００２％とした。逆に、１％超含有すると、炭窒化物の析出が多くなり加工性および耐遅れ破壊性低下が生じるため、上限値を１％とした。
【００４６】
更にＺｒは、鋼板の強度上昇及び細粒化に有効な元素である。しかし、０．００２％未満では析出物の個数が低下するために、下限値を０．００２％とした。逆に、１％超では粗大析出または昇出物が生成するために加工性および耐遅れ破壊性が低下する。このため、上限値を１％とした。
【００４７】
次に、Ｃｒ、Ｍｏ、Ｗは炭化物形成元素及び焼戻軟化抵抗元素であり、強度及び耐遅れ破壊性を改善するために必要な元素である。
更にＣｒは、鋼板の強度上昇に有効な元素である。しかし、０．００５％未満ではこれらの効果が得られないため、下限値を０．００５％とした。逆に、５％超含有すると加工性低下が生じるため、上限値を５％とした。
【００４８】
更にＭｏは、鋼板の焼入れ性を高め連続焼鈍設備で安定してマルテンサイトを得るために有効な元素であるだけでなく、粒界を強化して水素脆性の発生を抑制する効果がある。しかし、０．００５％未満ではこれらの効果が得られないため、下限値を０．００５％とした。また、５％超ではこれらの効果が飽和するため、上限値を５％とした。
【００４９】
更にＷは、鋼板の強度上昇に有効な元素である。しかし、０．００５％未満ではこれらの効果が得られないため、下限値を０．００５％とした。逆に、５％超含有すると加工性低下が生じるため、上限値を５％とした。
【００５０】
次に、Ｃｕは、強化に有効である上、自信の微細析出は遅れ破壊の向上にも寄与するため、０．００５％以上の添加とした。また、過剰添加は加工性の劣化を招くことから、上限を２．０％とした。
【００５１】
次に、Ｎｉ、Ｃｏは焼入れ性を高める強化元素である。
更にＮｉは、Ｎｉ硫化物が水素侵入を抑制し遅れ破壊特性を向上させる硬化や、鋼板の焼入れ性を高めることにより鋼板の強度を確保する効果がある。しかし、０．００５％未満ではこれらの効果が得られないため下限値を０．００５％とした。逆に、２％超では加工性が悪くなるため、上限値を２％とした。
【００５２】
更にＣｏは、強化に有効であるため、０．００５％以上の添加とした。また、過剰添加は加工性の劣化を招くことから、上限を２．０％とした。
【００５３】
次にＢは、鋼板の強度上昇に有効な元素である。しかし、０．０００２％未満ではこれらの効果が得られないため、下限値を０．０００２％とした。逆に、０．１％超含有すると熱間加工性が劣化するため、上限値を０．１％とした。
【００５４】
次にＲＥＭ、Ｃａ、Ｙは、介在物の形態制御に有効で、耐遅れ破壊性に寄与することから、０．０００５％以上の添加とした。一方、過剰添加は熱間加工性を劣化させるため、０．０１％以下の添加とした。
【００５５】
次に製造方法について説明する。
まず、規定した成分を有する鋳片を熱間圧延する。この際、フェライト粒にひずみが過度に加わり加工性が低下するのを防ぐためにＡｒ3 以上で仕上圧延を実施する。また、高温すぎても焼鈍後の再結晶粒径およびＭｇの複合析出または昇出物が必要以上に粗大化するため、９４０℃以下が望ましい。巻き取り温度については、高温にすれば再結晶や粒成長が促進され、加工性の向上が望まれるが、熱間圧延時に発生するスケール生成も促進され酸洗性が低下するので、８００℃以下とする。一方で低温になりすぎると硬化するため、冷間圧延時での負荷が高くなる。このため、５００℃以上とする。
【００５６】
酸洗後の冷間圧延は、圧下率が低いと鋼板の形状矯正が難しくなるため下限値を３０％とする。また、８０％を超える圧下率で圧延すると、鋼板のエッジ部に割れの発生及び形状の乱れのため上限値を８０％とする。連続焼鈍温度は低すぎると未再結晶の状態になり硬質化し、逆に高すぎると粒が粗大化しプレス時に肌荒れを起こす場合があるという問題点があるので、６００℃以上９５０℃以下とする。焼鈍は連続焼鈍または箱焼鈍設備を用いて行う。
【００５７】
また必要に応じて、焼鈍後、２００〜７００℃の温度域で１分から１０時間保持して、その後冷却しても良い。この熱処理により、合金炭化物または窒化物（例えばＶ、Ｃｒ、Ｍｏ、Ｗ含有の炭窒化物）を析出させ、これらが新たな水素のトラップサイトとして働き、より耐遅れ破壊性が高まる。条件が、低温短時間になると十分な析出が起こらず、高温長時間になると析出物が粗大化してトラップサイトとして機能しなくなることから、本範囲とした。
【００５８】
前記の鋳片は、鋳造速度が早いとＭｇ化合物が過度に微細化し、鋳造速度が遅いとＭｇ化合物が粗大化しかつ粒子数も少なくなりＭｇ化合物の遅れ破壊制御が十分に発揮できない場合が有る。鋳片の鋳造速度としては、０．０５ｍ／分〜２０．０ｍ／分が望ましい。更に、Ｍｇ化合物の遅れ破壊性向上効果を安定的に利用するには１．０ｍ／分〜３．０ｍ／分が好ましい。
【００５９】
なお、本鋼板は熱延鋼板、冷延鋼板、めっき鋼板のいずれでもかまわない。更にめっきは通常の亜鉛めっき、アルミめっき等のいずれでもかまわない。めっきは溶融めっき及び電気めっきのいずれでも良く、更にめっき後に合金化熱処理を施してもかまわないし、複層めっきでもかまわない。また、めっきを施さない鋼板上やめっき鋼板上にフィルムラミネート処理をした鋼板も本発明を逸脱するものではない。
【００６０】
さらに、本発明の高強度薄鋼板（例えば７８０ＭＰａ以上の鋼板）を形成加工した自動車用強度部品（例えばバンパーやドアインパクトビーム等の補強部材）においても、十分な材質特性（強度、剛性等）を表し、衝撃吸収性や耐遅れ破壊性も良好であった。
【００６１】
【実施例】
次に本発明を実施例に基づいて説明する．
表１に示す成分の鋼を溶製し、常法に従い連続鋳造でスラブとした。符号Ａ〜Ｊが本発明に従った成分の鋼で符号Ｋ〜Ｍは成分が逸脱するものである。これらの鋼を加熱炉中で１１６０℃〜１２５０℃の温度で加熱し、８７０℃〜９００℃の仕上げ温度で熱間圧延を行い、６５０℃〜７５０℃にて巻き取る。これに続いて符号Ｈ以外は、酸洗後に冷間圧延を行い、次いで再結晶焼鈍を行い、その後０．４％の調質圧延をして冷延鋼板となした。また、符号Ｉ、Ｊは更に目付量を片面５０ｇ／ｍ2 施した合金化溶融亜鉛めっき鋼板とし、Ｊについては加えてフィルムラミネート処理をした。表２に鋼板の製造方法と材質特性を示す。
【００６２】
表３に、鋼板の耐遅れ破壊性の評価を示した。評価方法は、８０mm×３０mmの短冊試験片を曲げ加工し、表面に耐水性の歪みゲージを装着した後で０．５mol/l の硫酸中に漬け、電流によって電解して水素を侵入させる方法である。その後割れの発生を評価した。曲げ加工の半径は５ｍｍ、１０ｍｍ、１５ｍｍとし、与える応力はそれぞれ６０ＭＰａと９０ＭＰａとした。
【００６３】
表２、表３に示すように、本発明例である符号１、２、３、５、７〜１２では、自動車の補強部品に適用するに充分な引張り強度と延性を示しており、割れ発生までの時間も長く耐遅れ破壊性に優れている。これらに対して比較例である符号４、６、１３〜１５では、成分、焼鈍温度のいずれかが本発明範囲から逸脱している。４、６は式（Ａ）の値が本範囲から逸脱しており、割れ発生までの時間が短くなる。１３〜１５は成分範囲が逸脱しており、水素のトラップサイトとなる晶出物や析出物の個数が少ないか、反対に水素をトラップしすぎるために割れ発生までの時間が短くなり、本発明での耐遅れ破壊性での差は明らかである。
【００６４】
【表１】
【表２】
【表３】
【発明の効果】
以上に説明した通り、本発明による高強度薄鋼板では、有効に水素のトラップサイトであるＭｇの化合物または複合晶出・析出物を効果的に分散させ、延性および成形加工後の耐遅れ破壊性を両立させる事ができる。本発明の高強度薄鋼板を形成加工した自動車用強度部品（例えばバンパーやドアインパクトビーム等の補強部材）においても、十分な材質特性を表し、衝撃吸収性や耐遅れ破壊性も良好であった。
【図面の簡単な説明】
【図１】式（Ａ）と遅れ破壊時間の関係を示すグラフである。
【図２】式（Ａ）と残留オーステナイトの関係を示すグラフである。
【図３】式（Ａ）とＭｇ量の関係を示すグラフである。
【図４】式（Ａ）と密度の関係を示すグラフである。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a high-strength thin steel sheet that suppresses cracking and delayed fracture, which is a problem particularly in a high-strength thin steel sheet, a method for producing the same, and a strength component for automobiles produced by them.
[0002]
[Prior art]
Conventionally, high-strength steel is often used for applications such as bolts, PC steel wires, and line pipes, and it is known that delayed fracture occurs due to hydrogen intrusion into the steel when the strength is 780 MPa or more. . In contrast, (1) the thin steel sheet is thin, so that hydrogen is released in a short time even if it penetrates, and (2) there is almost no use of a steel sheet of 780 MPa or more in terms of workability. It can be said that the awareness of problems with delayed destruction was low.
[0003]
However, recently, due to the necessity of reducing the weight of automobiles and improving collision safety, ultra high strength steel plates of 780 MPa or more are applied to stamping, pipe forming, bending, and end faces for reinforcing materials such as bumpers and impact beams and seat rails. The number of cases where processing and hole expansion processing are applied to use is increasing rapidly. Therefore, there is an urgent need to develop ultra high strength thin steel sheets with delayed fracture resistance.
[0004]
Until now, most technologies for improving delayed fracture resistance have been developed for steel materials that are often used as products such as bolts, strips, and thick plates, and are used below the yield strength or yield stress.
[0005]
For example, in steel bars and bolt steels, development was centered on tempered martensite, and “New Developments in Elucidation of Delayed Fracture” (Japan Iron and Steel Institute, Material Structure and Properties Subcommittee, High Strength Steel Delayed Fracture Study Group) In January 1997), it was reported that additive elements exhibiting temper softening resistance such as Cr, Mo and V are effective in improving delayed fracture resistance. This is a technique for precipitating alloy carbides and using them as hydrogen trap sites to shift the delayed fracture mode from grain boundaries to intragranular fracture. However, these steels have a C content of 0.4% or more and contain a large amount of alloy elements, so the workability and weldability required for thin steel sheets are inferior, and the precipitation of alloy carbide takes several hours or more. Therefore, there is a problem in manufacturability. JP-A-11-293383 describes that an oxide mainly composed of Ti and Mg is effective in preventing hydrogen defects. However, this is intended for thick steel plates, especially considering delayed fracture after welding with high heat input, but undergoes high forming work required for thin steel plates and burrs associated with end face processing. No consideration is given to the effects on delayed fracture phenomena such as occurrence. Furthermore, there is no consideration for workability, which is a basic characteristic of thin steel plates.
[0006]
On the other hand, regarding delayed fracture of thin steel sheets (for example, Yamazaki et al. CAMP-ISIJ vol.5 p1839 to 1842 (1992)), it has been reported about the promotion of delayed fracture due to work-induced transformation of the amount of retained austenite. This is in consideration of the forming process of a thin steel sheet, but describes the regulation of the amount of retained austenite which does not deteriorate the delayed fracture resistance. That is, it relates to a high-strength thin steel sheet having a specific structure, and is not a fundamental countermeasure for improving delayed fracture resistance.
[0007]
[Problems to be solved by the invention]
As mentioned above, taking into account the use environment of thin steel sheets and manufacturability with current equipment, development examples that take measures against delayed fracture such as forming before use while securing the basic formability are as follows: rare.
[0008]
[Means for Solving the Problems]
From the background as described above, the present inventors have come to find a method for fundamentally improving delayed fracture resistance by sufficiently considering the use environment in thin steel sheets and the manufacturing method using current equipment. That is, it has been found that by controlling the form by forming a Mg compound or a composite compound, the delayed fracture resistance after forming is improved without deteriorating the formability of the high-strength thin steel sheet. . And the effective manufacturing method of the said high intensity | strength thin steel plate using the present manufacturing equipment (Hot rolling, continuous annealing, hot dip galvanization, electroplating equipment, etc.) was discovered. Details are as follows.
[0009]
(1) By weight%
C: 0.05% to 0.3%,
Si: 3.0% or less,
Mn: 0.01 to 3.0%,
P: 0.02% or less,
S: 0.02% or less,
Al: 0.01% to 3.0%,
N: 0.01% or less,
Mg: 0.0002% to 0.01%
In which the balance is steel composed of iron and inevitable impurities, the retained austenite is 7% or less by volume in the structure of the steel sheet, Mg oxide, Mg sulfide, Mg Any one or more of the composite crystallized product and the composite precipitate of Mg are
Average particle diameter d: 0.01 to 5.0 μm
Density ρ: 100 to 100,000 per square mm Distribution: ratio of standard deviation σ from average particle diameter to average particle diameter d σ / d ≦ 1.0
And the average particle diameter d, the density ρ, the volume fraction Vγ ( % ) of retained austenite, and the tensile strength TS ( MPa ) satisfy the formula (A). High-strength thin steel sheet with excellent delayed fracture resistance after processing.
Formula (A): 1000 (Vγ−0.1) −5.5 + α (Mg−40) 2 −50 (d−0.2) 2 + 1.1lnρ + 700 (TS−680) −0.9 ≧ 10
*) Α = −0.005 (Mg ≦ 40), α = −0.002 (Mg> 40)
Vγ: Volume fraction of retained austenite (%)
Mg: Mg amount (ppm by weight)
d: Average particle diameter (μm)
ρ: density (pieces / mm 2 )
TS: Tensile strength (MPa)
Conditions of formula (A): [1] 1000 (V γ -0.1) is 1000 (Vγ-0.1) When -5.5 ≧ 10 -5.5 = 10
[2] 2 ≦ Mg ≦ 100ppm
[3] When 0.01 ≦ d ≦ 5.0 μm and (d−0.2) 2 ≦ 0.2, (d−0.2) 2 = 0.2
[4] 100 ≦ ρ ≦ 100,000 pieces / mm 2
[5] 780 MPa ≦ TS
[0010]
(2) Furthermore, in weight%,
V: 0.005 to 1%,
Ti: 0.002 to 1%,
Nb: 0.002 to 1%
Zr: 0.002 to 1%,
The high-strength thin steel sheet having excellent delayed fracture resistance after forming according to (1), characterized in that the steel contains one or more of the following, the balance being iron and inevitable impurities.
[0011]
(3) Furthermore, in weight%,
Cr: 0.005 to 5%,
Mo: 0.005 to 5%,
W: 0.005 to 5%
Of one or comprise two or more, the balance being a steel comprising iron and unavoidable impurities (1) or high excellent delayed fracture resistance after molding according to (2) Strength thin steel plate.
[0012]
(4) Furthermore, in weight%,
Cu: 0.005 to 2.0%
A high-strength thin steel sheet excellent in delayed fracture resistance after forming according to any one of (1) to (3), wherein the balance is steel made of iron and inevitable impurities.
[0013]
(5) Furthermore, in weight%,
Ni: 0.005 to 2.0%,
Co: 0.005 to 2.0%
(1) to (4) characterized in that it contains one or more of the following, and the balance is steel composed of iron and unavoidable impurities. Strength thin steel plate.
[0014]
(6) Furthermore, in weight%,
B: 0.0002 to 0.1%
A high-strength thin steel sheet excellent in delayed fracture resistance after forming according to any one of (1) to (5), wherein the balance is steel made of iron and inevitable impurities.
[0015]
(7) Furthermore, in weight%,
REM: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
Y: 0.0005 to 0.01%
(1) to (6) characterized in that it has one or more of the following, and the balance is steel composed of iron and unavoidable impurities. Strength thin steel plate.
[0016]
(8) A slab having the composition according to any one of (1) to (7) is manufactured, hot-rolled at a finishing temperature of Ar 3 point or higher, scraped at 500 to 800 ° C., and then normal After pickling, after cold rolling with a rolling reduction of 30 to 80%, soaking at 600 ° C. to 950 ° C. in the annealing step, recrystallization annealing is performed, and then lag resistance after forming processing is subjected to temper rolling Manufacturing method of high strength thin steel sheet with excellent destructibility.
[0017]
(9) In the production method according to (8), a high-strength thin steel sheet excellent in delayed fracture resistance after forming, characterized by holding in a temperature range of 200 to 700 ° C. for 1 minute to 10 hours after annealing. Manufacturing method.
[0018]
(10) A high strength thin steel sheet having excellent delayed fracture resistance after forming, wherein the high strength steel sheet according to (1) to (7) is a hot rolled steel sheet or a cold rolled steel sheet.
[0019]
(11) A high-strength thin steel sheet excellent in delayed fracture resistance after forming according to any one of (1) to (7), wherein the steel sheet is subjected to a galvanized surface treatment.
[0020]
(12) A high-strength thin steel sheet excellent in delayed fracture resistance after forming, in which the steel sheet described in (9) or (10) is further subjected to film lamination.
[0021]
(13) An automotive strength component made of the high-strength thin steel sheet according to (1) to (7).
[0022]
DETAILED DESCRIPTION OF THE INVENTION
In tempered martensitic steel, it is considered that delayed fracture is caused by the accumulation of hydrogen at the prior austenite grain boundaries and the like, resulting in the formation of voids and the like, which cause fracture. Therefore, if the hydrogen trap sites are dispersed evenly and finely and hydrogen is trapped there, the concentration of diffusible hydrogen is lowered and the susceptibility to delayed fracture is lowered. As described in the above-mentioned JP-A-11-293383, it has been found that delayed fracture resistance due to hydrogen is improved by controlling the oxide dispersion form in a thick steel plate to which Mg and Ti are added in combination. However, because it undergoes forming processing like a thin steel plate, high residual stress is generated, and if there are burrs on the processing end face, the delayed fracture resistance naturally deteriorates. I cannot supplement. Thus, there are few studies on delayed fracture characteristics in consideration of the usage pattern of thin steel sheets, and it cannot be solved only by controlling the oxide morphology of Mg or Ti. Further, there is a concern that the fine dispersion of the trap site deteriorates the ductility, which is a basic characteristic of the thin steel plate.
[0023]
In light of the above-mentioned background, the present inventors ensured and improved delayed fracture resistance even after forming, in addition to various crystallized materials and precipitates, The effects of strength and organization were examined. As a result, we have found a technique for improving and ensuring delayed fracture resistance even under high residual stress and end face burrs in the environment where thin steel sheets are used. That is,
(1) Control of dispersion form of Mg-containing oxides or sulfides and compounds crystallized or precipitated with them.
(2) Amount of retained austenite in the microstructure of the steel sheet.
(3) Strength of steel sheet.
By controlling each of these, it is possible to effectively disperse Mg compounds or composite crystallization / precipitates that are hydrogen trap sites, and to achieve both ductility and delayed fracture resistance after forming. Formula (A) was defined as a condition for satisfying this. It is to be noted that the oxide or sulfide containing Mg of (1) and the compound crystallized or precipitated with them are more delayed when they are in the grains (excluding the phase interface of the microstructure such as the prior austenite grain boundaries). Effective in improving fracture characteristics.
[0024]
In order to obtain (1), (2) and (3) of the present invention described above, the production conditions are specified, and crystallized substances and precipitates such as oxides, nitrides and sulfides of various elements are trapping sites of hydrogen. It is also effective to control the form that can be.
[0025]
By satisfying the formula (A), the delayed fracture resistance of the high-strength thin steel sheet can be sufficiently secured even after the forming process.
This is because the interaction between the dislocations and residual stress fields introduced by the forming process of thin steel sheets and the particles that become trap sites is the same as the dislocations and residual stresses introduced during hot rolling and cooling after welding on thick steel sheets. It is thought that it originates in a difference and the difference in the heat processing method of a thin steel plate and a thick steel plate.
The above (1) and (2) are limited as follows.
[0026]
Residual austenite amount: Residual austenite increases delayed fracture susceptibility when it becomes martensite due to processing-induced transformation, so the upper limit was made 7% or less in volume ratio.
[0027]
Average particle diameter: The average particle diameter was limited to 0.01 μm to 5.0 μm. The hydrogen trap site needs to have a certain size, and a large amount of fine particles is not preferable from the viewpoint of ensuring the ductility of the thin steel sheet and is difficult to manufacture. Therefore, the lower limit was set to 0.01 μm. Further, since the coarse particles do not act as trap sites and can be a starting point of destruction, the upper limit is set to 5.0 μm.
[0028]
Density: The particle density was 100 to 100,000 / mm 2 . Low particle density means that the number of trap sites is small, and delayed fracture resistance after processing cannot be secured, so the lower limit was set to 100 particles / mm 2 . In the case of high density, the ductility and molding processability deteriorate and the effect of improving delayed fracture resistance is saturated, so the upper limit was set to 100,000 pieces / mm 2 .
[0029]
Distribution: The particle distribution is such that the ratio σ / d of the standard deviation σ from the average particle diameter to the average particle diameter d satisfies σ / d ≦ 1.0 . σ / d> 1.0 means that the particle distribution covers a wide range, and the effect of improving delayed fracture resistance is smaller than that of the same average particle size, which leads to ductile deterioration and an increase in the number of fracture starting points. The upper limit was made 1.0 or less.
[0030]
Here, measurement of particles containing an Mg compound will be described. The particle is measured by using a thin film or sample of an extraction replica with a scanning or transmission electron microscope at a magnification of 5,000 to 100,000, and measuring at least 30 fields. The particle diameter is evaluated by the equivalent circle diameter by image analysis. Moreover, when calculating | requiring a density, a composite precipitation or a crystallization thing is counted as one piece. Composition analysis was performed by using EDX and EELS, and structural analysis was performed by analyzing a Diffraction pattern. Each composite compound is a compound (carbide, nitride, oxide, sulfide, etc.) containing an alloy addition element (for example, Ti, Nb, V, Cr, Mo, REM, Ca, etc.) in addition to Mg.
[0031]
The present invention is described in further detail below.
In the present invention, a high-strength thin steel plate is described, but mainly a steel plate having a tensile strength of 780 MPa or more and a plate thickness of 0.5 mm to 4.0 mm.
Next, in formula (A), the volume fraction of retained austenite, average particle diameter, density, Mg amount, and tensile strength were set from FIG. 1 in addition to being cited as factors of delayed fracture resistance. When the left side of the formula (A) is a function f (Vγ, Mg, d, ρ, TS), the value of f (Vγ, Mg, d, ρ, TS) is 10 or more, and the delayed fracture resistance is remarkably improved. . Furthermore, the influence with respect to delayed fracture resistance about each variable is shown in FIGS. ○ in the figure indicates that the delayed fracture resistance is good, and x indicates that it is defective.
[0032]
FIG. 2 is a correlation graph between f (Vγ) and residual austenite volume fraction Vγ. The conditions of the graph are Mg: 30 ppm, average particle size: 0.4 μm, density: 1500 particles / mm 2 , and tensile strength: 1480 MPa. When Vγ is high, delayed fracture resistance deteriorates, but the steel according to the invention having a high f (Vγ) value exhibits good delayed fracture resistance when Vγ is 7% or less. Further, even when Vγ is 7% or less, the comparative steel of × is out of the range because Mg, particle size, and density are f (Vγ) <10, and the delayed fracture resistance deteriorated.
[0033]
FIG. 3 is a correlation graph of f (Mg) and Mg addition amount. The conditions of the graph are: volume ratio of retained austenite: 3.0%, average particle size: 0.4 μm, density: 1500 particles / mm 2 , and tensile strength: 1480 MPa. There is a region where the delayed fracture resistance is particularly good when the Mg addition amount is 20 to 70 ppm. Further, Mg with 100 ppm or less and x indicates that the retained austenite, particle size, and density are out of range, and therefore f (Mg) <10 and delayed fracture resistance deteriorated.
[0034]
FIG. 4 is a correlation graph of f (ρ) and the density of crystallized substances and precipitates. The conditions of the graph are the volume ratio of retained austenite: 3.0%, Mg: 30 ppm, and tensile strength: 1380 MPa. If the density is low, the delayed fracture resistance is poor. Moreover, although the density ρ is within the range of the claims, the case of x is because the retained austenite, Mg, and the particle size are out of the range, so f (ρ) <10 and the delayed fracture resistance deteriorated.
From the above, it was assumed that the delayed fracture resistance was excellent if the formula (A) was satisfied.
[0035]
Next, the reason for limiting the chemical composition of steel in the present invention will be described.
C is an element that can increase the strength of the steel sheet. In particular, it is an essential element for generating a hard phase such as martensite and austenite and increasing the strength, and 0.05% or more is necessary to obtain a strength of 780 MPa or more. In order to increase the cementite that is the starting point of fracture, hydrogen embrittlement is likely to occur. Therefore, the upper limit was made 0.3%.
[0036]
Si is a substitutional solid solution strengthening element that greatly hardens the material, and is effective in increasing the strength of the steel sheet, and is an element that suppresses cementite precipitation. The scale removal is costly and economically disadvantageous, so 3.0% is made the upper limit. Further, if the amount added is large, the plating property is deteriorated, so 0.6% or less is desirable for improving the plating property.
[0037]
Mn is an element effective for increasing the strength of the steel sheet. However, since this effect cannot be obtained if the content is less than 0.01%, the lower limit is set to 0.01%. On the contrary, if the amount is large, not only co-segregation with P and S is promoted, but also workability may be deteriorated, so 3.0% is made the upper limit.
[0038]
P is an element that promotes grain boundary fracture due to grain boundary segregation, and a lower value is desirable, but an extreme decrease is not preferable in terms of manufacturing cost. Moreover, since it is an element which degrades corrosion resistance, the upper limit is made 0.02%.
[0039]
S is an element that promotes hydrogen absorption in a corrosive environment, and a lower value is desirable. However, extreme reduction is not preferable in terms of manufacturing cost. In particular, in order to improve workability, a lower value is desirable and the upper limit is set to 0.02%.
[0040]
Al is added in an amount of 0.01% or more for deoxidation. Increasing the amount increases inclusions such as alumina, and deteriorates workability and weldability, so 3.0% is the upper limit. To do. It should be noted that addition of 0.2% or more is desirable because it has an effect of suppressing the formation of retained austenite.
[0041]
N is better because it contributes to workability deterioration and blowhole generation during welding. If it exceeds 0.01%, workability deteriorates, so 0.01% is made the upper limit.
[0042]
Mg is not only effective for improving delayed fracture resistance of its own compound, but also forms composite precipitates or composite crystallization products with other elements, and contributes to improving delayed fracture resistance of these forms. Since it is an element necessary for such control, it was made 0.0002% or more. However, if it exceeds 0.01%, coarse oxides and sulfides are generated, which is not effective for shape control, and the formability, which is a basic required characteristic of a thin steel sheet, is reduced. %.
[0043]
Next, V, Ti, Nb, and Zr are strong carbide generating elements, and are necessary for generating precipitates and inclusions to improve strength and delayed fracture resistance.
Furthermore, V is an element effective for increasing the strength of the steel sheet and reducing the grain size. However, since this effect cannot be obtained at less than 0.005%, the lower limit is set to 0.005%. On the other hand, when the content exceeds 1%, the precipitation of carbonitrides becomes remarkable, and the ductility decreases remarkably. For this reason, the upper limit is set to 1%.
[0044]
Further, Ti is an element effective for increasing the strength of the steel sheet and reducing the grain size. However, if it is less than 0.002%, the number of precipitates decreases, so the lower limit was made 0.002%. On the other hand, if it exceeds 1%, coarse precipitates or ascendants are formed, so that workability and delayed fracture resistance deteriorate. For this reason, the upper limit is set to 1%.
[0045]
Further, Nb is an element effective for increasing the strength and refining of the steel sheet. However, since these effects cannot be obtained if the content is less than 0.002%, the lower limit is set to 0.002%. On the other hand, when the content exceeds 1%, precipitation of carbonitride increases, resulting in deterioration of workability and delayed fracture resistance. Therefore, the upper limit is set to 1%.
[0046]
Furthermore, Zr is an element effective for increasing the strength and refining of the steel sheet. However, if it is less than 0.002%, the number of precipitates decreases, so the lower limit was made 0.002%. On the other hand, if it exceeds 1%, coarse precipitates or ascendants are formed, so that workability and delayed fracture resistance deteriorate. For this reason, the upper limit is set to 1%.
[0047]
Next, Cr, Mo, and W are carbide forming elements and temper softening resistance elements, and are necessary elements for improving strength and delayed fracture resistance.
Furthermore, Cr is an effective element for increasing the strength of the steel sheet. However, since these effects cannot be obtained at less than 0.005%, the lower limit is set to 0.005%. On the other hand, if the content exceeds 5%, the workability deteriorates, so the upper limit was made 5%.
[0048]
Furthermore, Mo is not only an effective element for improving the hardenability of the steel sheet and stably obtaining martensite in a continuous annealing facility, but also has an effect of strengthening grain boundaries and suppressing the occurrence of hydrogen embrittlement. However, since these effects cannot be obtained at less than 0.005%, the lower limit is set to 0.005%. Further, if the content exceeds 5%, these effects are saturated, so the upper limit is set to 5%.
[0049]
Furthermore, W is an element effective for increasing the strength of the steel sheet. However, since these effects cannot be obtained at less than 0.005%, the lower limit is set to 0.005%. On the other hand, if the content exceeds 5%, the workability deteriorates, so the upper limit was made 5%.
[0050]
Next, Cu is effective for strengthening, and the fine precipitation of confidence contributes to the improvement of delayed fracture, so 0.005% or more was added. Moreover, since excessive addition causes deterioration of workability, the upper limit was made 2.0%.
[0051]
Next, Ni and Co are strengthening elements that enhance hardenability.
Furthermore, Ni has the effect of ensuring the strength of the steel sheet by hardening the Ni sulfide to suppress hydrogen intrusion and improving delayed fracture characteristics, and enhancing the hardenability of the steel sheet. However, if the content is less than 0.005%, these effects cannot be obtained, so the lower limit is set to 0.005%. On the contrary, if it exceeds 2%, the workability deteriorates, so the upper limit was set to 2%.
[0052]
Furthermore, Co is effective for strengthening, so 0.005% or more was added. Moreover, since excessive addition causes deterioration of workability, the upper limit was made 2.0%.
[0053]
Next, B is an element effective for increasing the strength of the steel sheet. However, since these effects cannot be obtained if the content is less than 0.0002%, the lower limit is set to 0.0002%. On the other hand, if the content exceeds 0.1%, the hot workability deteriorates, so the upper limit was made 0.1%.
[0054]
Next, REM, Ca, and Y are effective for controlling the form of inclusions and contribute to delayed fracture resistance, so 0.0005% or more was added. On the other hand, excessive addition deteriorates hot workability, so 0.01% or less was added.
[0055]
Next, a manufacturing method will be described.
First, a slab having a specified component is hot-rolled. At this time, finish rolling is performed with Ar 3 or more in order to prevent distortion from being excessively applied to the ferrite grains and deterioration of workability. Further, even if the temperature is too high, the recrystallized grain size after annealing and the composite precipitation or ascending product of Mg are coarsened more than necessary, so that the temperature is preferably 940 ° C. or lower. With regard to the coiling temperature, recrystallization and grain growth are promoted at a high temperature, and improvement in workability is desired. However, scale formation that occurs during hot rolling is also promoted and pickling properties are reduced. And On the other hand, since it hardens | cures when it becomes low temperature, the load at the time of cold rolling becomes high. For this reason, it shall be 500 degreeC or more.
[0056]
In cold rolling after pickling, if the rolling reduction is low, it becomes difficult to correct the shape of the steel sheet, so the lower limit is set to 30%. Further, when rolling at a rolling reduction exceeding 80%, the upper limit is set to 80% because of the occurrence of cracks and the disorder of the shape of the edge of the steel sheet. If the continuous annealing temperature is too low, it becomes a non-recrystallized state and becomes hard, and conversely if too high, there is a problem that the grains become coarse and rough skin may occur during pressing. Annealing is performed using continuous annealing or box annealing equipment.
[0057]
If necessary, after annealing, it may be kept in a temperature range of 200 to 700 ° C. for 1 minute to 10 hours and then cooled. By this heat treatment, alloy carbides or nitrides (for example, carbonitrides containing V, Cr, Mo, W) are precipitated, and these act as new hydrogen trap sites, further increasing delayed fracture resistance. If the conditions are low temperature for a short time, sufficient precipitation does not occur, and if the temperature is high for a long time, the precipitate becomes coarse and does not function as a trap site.
[0058]
In the above slab, when the casting speed is high, the Mg compound becomes excessively fine, and when the casting speed is low, the Mg compound becomes coarse and the number of particles decreases, and the delayed fracture control of the Mg compound may not be fully exhibited. The casting speed of the slab is preferably 0.05 m / min to 20.0 m / min. Furthermore, 1.0 m / min to 3.0 m / min is preferable in order to stably utilize the delayed fracture property improving effect of the Mg compound.
[0059]
In addition, this steel plate may be a hot rolled steel plate, a cold rolled steel plate, or a plated steel plate. Furthermore, the plating may be any of normal galvanizing and aluminum plating. Plating may be either hot dipping or electroplating, and may be further subjected to alloying heat treatment after plating, or may be multilayer plating. Moreover, the steel plate which carried out the film lamination process on the steel plate which does not give plating, or a plated steel plate does not deviate from this invention.
[0060]
Furthermore, sufficient strength characteristics (strength, rigidity, etc.) are also obtained in automotive strength parts (for example, reinforcing members such as bumpers and door impact beams) formed and processed the high-strength thin steel sheets (for example, steel sheets of 780 MPa or more) of the present invention. The impact absorption and delayed fracture resistance were also good.
[0061]
【Example】
Next, the present invention will be described based on examples.
Steels having the components shown in Table 1 were melted and slabs were obtained by continuous casting according to a conventional method. The symbols A to J are steels of the components according to the present invention, and the symbols K to M are components that deviate. These steels are heated in a heating furnace at a temperature of 1160 ° C to 1250 ° C, hot rolled at a finishing temperature of 870 ° C to 900 ° C, and wound up at 650 ° C to 750 ° C. Subsequently, except for the symbol H, cold rolling was performed after pickling, followed by recrystallization annealing, followed by 0.4% temper rolling to obtain a cold rolled steel sheet. Further, symbols I and J were alloyed hot-dip galvanized steel sheets having a basis weight of 50 g / m 2 on one side, and J was additionally subjected to film lamination. Table 2 shows the steel sheet manufacturing method and material properties.
[0062]
Table 3 shows the evaluation of delayed fracture resistance of the steel sheet. The evaluation method is a method of bending a strip test piece of 80 mm x 30 mm, mounting a water-resistant strain gauge on the surface, then immersing it in 0.5 mol / l sulfuric acid, and injecting hydrogen by electrolysis with current. is there. Thereafter, the occurrence of cracks was evaluated. The bending radius was 5 mm, 10 mm, and 15 mm, and the applied stress was 60 MPa and 90 MPa, respectively.
[0063]
As shown in Tables 2 and 3,
[0064]
[Table 1]
[Table 2]
[Table 3]
【The invention's effect】
As explained above, the high strength thin steel sheet according to the present invention effectively disperses Mg compounds or composite crystallization / precipitates which are hydrogen trap sites, and has ductility and delayed fracture resistance after forming. Can be made compatible. The automotive strength parts (for example, reinforcing members such as bumpers and door impact beams) formed and processed the high-strength thin steel sheet of the present invention also showed sufficient material properties, and good shock absorption and delayed fracture resistance. .
[Brief description of the drawings]
FIG. 1 is a graph showing the relationship between formula (A) and delayed fracture time.
FIG. 2 is a graph showing the relationship between formula (A) and retained austenite.
FIG. 3 is a graph showing the relationship between the formula (A) and the amount of Mg.
FIG. 4 is a graph showing the relationship between equation (A) and density.
Claims (13)
Ｃ ：０．０５％〜０．３％、
Ｓｉ：３．０％以下、
Ｍｎ：０．０１〜３．０％、
Ｐ ：０．０２％以下、
Ｓ ：０．０２％以下、
Ａｌ：０．０１％〜３．０％、
Ｎ ：０．０１％以下、
Ｍｇ：０．０００２％〜０．０１％
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とし、鋼板の組織中に残留オーステナイトが体積率で７％以下であり、Ｍｇの酸化物、Ｍｇの硫化物、Ｍｇの複合晶出物およびＭｇの複合析出物のいずれか１種以上が、
平均粒子径ｄ：０．０１〜５．０μｍ
密度ρ：１平方ｍｍあたり１００〜１０００００個
分布：平均粒子径からの標準偏差σと平均粒子径ｄの比σ / ｄ≦１．０
を満たす分布形態を有し、さらに、前記平均粒子径ｄ、前記密度ρ、残留オーステナイトの体積率Ｖγ ( ％ ) 、引張強度ＴＳ ( ＭＰａ ) が式（Ａ）を満たすことを特徴とする、成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。
式（Ａ）：１０００（Ｖγ−０．１）−５．５＋α（Ｍｇ−４０)２−５０（ｄ−０．２)２＋１．１ｌｎρ＋７００（ＴＳ−６８０）−０．９≧１０
＊）α＝−０．００５（Ｍｇ≦４０）、α＝−０．００２（Ｍｇ＞４０）
Ｖγ：残留オーステナイト体積率（％）
Ｍｇ：Ｍｇ量（重量ｐｐｍ）
ｄ ：平均粒子径（μｍ）
ρ ：密度（個／ｍｍ２）
ＴＳ：引張強度（ＭＰａ）
式（Ａ）の条件：[1]１０００（Ｖγ−０．１）−５．５≧１０の時は１０００（Ｖγ−０．１）−５．５＝１０
[2]２≦Ｍｇ≦１００ｐｐｍ
[3]０．０１≦ｄ≦５．０μｍ、（ｄ−０．２）２≦０．２の時は（ｄ−０．２）２ ＝０．２
[4]１００≦ρ≦１０００００個／ｍｍ２
[5]７８０ＭＰａ≦ＴＳ% By weight
C: 0.05% to 0.3%,
Si: 3.0% or less,
Mn: 0.01 to 3.0%,
P: 0.02% or less,
S: 0.02% or less,
Al: 0.01% to 3.0%,
N: 0.01% or less,
Mg: 0.0002% to 0.01%
In which the balance is steel composed of iron and inevitable impurities, the retained austenite is 7% or less by volume in the structure of the steel sheet, Mg oxide, Mg sulfide, Mg Any one or more of the composite crystallized product and the composite precipitate of Mg are
Average particle diameter d: 0.01 to 5.0 μm
Density ρ: 100 to 100,000 per square mm Distribution: ratio of standard deviation σ from average particle diameter to average particle diameter d σ / d ≦ 1.0
And the average particle diameter d, the density ρ, the volume fraction Vγ ( % ) of retained austenite, and the tensile strength TS ( MPa ) satisfy the formula (A). High-strength thin steel sheet with excellent delayed fracture resistance after processing.
Formula (A): 1000 (Vγ−0.1) −5.5 + α (Mg−40) 2 −50 (d−0.2) 2 + 1.1lnρ + 700 (TS−680) −0.9 ≧ 10
*) Α = −0.005 (Mg ≦ 40), α = −0.002 (Mg> 40)
Vγ: Volume fraction of retained austenite (%)
Mg: Mg amount (ppm by weight)
d: Average particle diameter (μm)
ρ: density (pieces / mm 2 )
TS: Tensile strength (MPa)
Conditions of formula (A): [1] 1000 (V γ -0.1) is 1000 (Vγ-0.1) When -5.5 ≧ 10 -5.5 = 10
[2] 2 ≦ Mg ≦ 100ppm
[3] When 0.01 ≦ d ≦ 5.0 μm and (d−0.2) 2 ≦ 0.2, (d−0.2) 2 = 0.2
[4] 100 ≦ ρ ≦ 100,000 pieces / mm 2
[5] 780 MPa ≦ TS
Ｖ ：０．００５〜１％、
Ｔｉ：０．００２〜１％、
Ｎｂ：０．００２〜１％、
Ｚｒ: ０．００２〜１％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
V: 0.005 to 1%,
Ti: 0.002 to 1%,
Nb: 0.002 to 1%
Zr: 0.002 to 1%
The high strength thin steel sheet having excellent delayed fracture resistance after forming according to claim 1, wherein the steel is one or more of the following, the balance being iron and inevitable impurities.
Ｃｒ：０．００５〜５％、
Ｍｏ：０．００５〜５％、
Ｗ ：０．００５〜５％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１または２に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
Cr: 0.005 to 5%,
Mo: 0.005 to 5%,
W: 0.005 to 5%
A high-strength thin film excellent in delayed fracture resistance after forming according to claim 1 or 2, wherein the steel is composed of one or more of the following, the balance being iron and inevitable impurities. steel sheet.
Ｃｕ：０．００５〜２．０％
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１〜３に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
Cu: 0.005 to 2.0%
The high-strength thin steel sheet excellent in delayed fracture resistance after forming according to claim 1, wherein the balance is steel made of iron and inevitable impurities.
Ｎｉ：０．００５〜２．０％、
Ｃｏ：０．００５〜２．０％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１から４に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
Ni: 0.005 to 2.0%,
Co: 0.005 to 2.0%
A high-strength thin film excellent in delayed fracture resistance after forming according to any one of claims 1 to 4, characterized in that it comprises one or more of the following, the balance being iron and inevitable impurities. steel sheet.
Ｂ ：０．０００２〜０．１％
を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１から５に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
B: 0.0002 to 0.1%
The high-strength thin steel sheet having excellent delayed fracture resistance after forming according to claim 1, wherein the balance is steel made of iron and inevitable impurities.
ＲＥＭ：０．０００５〜０．０１％、
Ｃａ：０．０００５〜０．０１％、
Ｙ ：０．０００５〜０．０１％
の１種または２種以上を含有し、残部が鉄および不可避的不純物からなる鋼であることを特徴とする請求項１から６に記載の成形加工後の耐遅れ破壊性に優れた高強度薄鋼板。Furthermore, in weight%
REM: 0.0005 to 0.01%,
Ca: 0.0005 to 0.01%,
Y: 0.0005 to 0.01%
A high-strength thin film excellent in delayed fracture resistance after forming according to claim 1, wherein the steel is composed of one or more of the following, the balance being iron and inevitable impurities. steel sheet.
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