JP2004204341A - Hot-dip galvanized steel sheet having ultrafine grain structure and excellent ductility and method for producing the same - Google Patents

Abstract

An object of the present invention is to provide a high-strength hot-dip galvanized steel sheet having a fine grain structure and excellent mechanical properties, particularly ductility.
SOLUTION: Among steel components, C, Si, Mn, Ni, Ti and Nb are contained in a range satisfying the following formulas (1), (2) and (3), respectively, and the balance is Fe and unavoidable impurities. A steel structure having a ferrite volume fraction of at least 70 vol%, a retained austenite volume fraction of at least 5 vol%, and an average crystal grain size of the ferrite of 3.5 μm or less.
637.5 + 4930 {Ti ^* + (48/93) · [% Nb]} ≧ A ₁ --- (1)
A ₃ ≤ 860 --- (2)
[% Mn] + [% Ni] ≧ 1.3 −−− (3)
^{However, Ti * = [% Ti]} - (48/32) · [% S] - (48/14) · [% N], A 1: predicted value of _{the A 1} transformation point obtained by calculation formula (℃) , A ₃ : Predicted value of A ₃ transformation point obtained by calculation formula (° C.)
[Selection diagram] Fig. 1

Description

【００１５】
２．上記１において、鋼板が、質量％でさらに、
Ｍｏ：１．０ ％以下および
Ｃｒ：１．０ ％以下
のうちから選んだ一種または二種を含有する組成になることを特徴とする、超微細粒組織を有し延性に優れる溶融亜鉛めっき鋼板。
【００１６】
３．上記１または２において、鋼板が、質量％でさらに、
Ｃａ， ＲＥＭおよびＢのうちから選んだ一種または二種以上を合計で０．００５ ％以下
を含有する組成になることを特徴とする、超微細粒組織を有し延性に優れる溶融亜鉛めっき鋼板。
【００１７】
４．質量％で、
Ｃ：０．０３〜０．１６％、
Ｓｉ：０．２ 〜２．０ ％、
Ｍｎ：１．０ 〜３．０ ％および／またはＮｉ：０．５ 〜３．０ ％、
Ｔｉ：０．２ ％以下および／またはＮｂ：０．２ ％以下、
Ａｌ：０．０１〜０．１ ％、
Ｐ：０．１ ％以下、
Ｓ：０．０２％以下および
Ｎ：０．００５ ％以下
で、かつＣ，Ｓｉ， Ｍｎ， Ｎｉ， ＴｉおよびＮｂが下記（１）， （２）， （３） 式をそれぞれ満足する範囲において含有し、残部はＦｅおよび不可避的不純物の組成になる鋼素材を、１２００℃以上に加熱したのち、熱間圧延し、ついで冷間圧延後、下記（６） 式で求められる温度Ａ_３ ℃以上、（Ａ_３ ＋３０） ℃以下で再結晶焼鈍を施し、ついで酸洗後、下記（５） 式で求められるＡ_１ ℃以上、（Ａ_１ ＋７０） ℃以下の温度範囲で５〜３０秒の熱処理を施し、引き続き溶融亜鉛めっき処理、あるいはさらに合金化処理を施すことを特徴とする、超微細粒組織を有し延性に優れる溶融亜鉛めっき鋼板の製造方法。

【００１８】
５．上記４において、鋼素材が、質量％でさらに、
Ｍｏ：１．０ ％以下および
Ｃｒ：１．０ ％以下
のうちから選んだ一種または二種を含有する組成になることを特徴とする、超微細粒組織を有し延性に優れる溶融亜鉛めっき鋼板の製造方法。
【００１９】
６．上記４または５において、鋼素材が、質量％でさらに、
Ｃａ， ＲＥＭおよびＢのうちから選んだ一種または二種以上を合計で０．００５ ％以下
を含有する組成になることを特徴とする、超微細粒組織を有し延性に優れる溶融亜鉛めっき鋼板の製造方法。
【００２０】
【発明の実施の形態】
以下、本発明を具体的に説明する。
まず、本発明において鋼素材の成分組成を上記の範囲に限定した理由について説明する。なお、成分に関する「％」表示は特に断らない限り質量％を意味するものとする。
Ｃ：０．０３〜０．１６％
Ｃは、安価な強化成分であるだけでなく、延性に寄与する残留オーステナイトを安定化させる上でも有用な元素である。しかしながら、含有量が０．０３％に満たないとその添加効果に乏しく、一方０．１６％を超えて含有させると溶接性が劣化するため、Ｃは０．０３〜０．１６％の範囲に限定した。
【００２１】
Ｓｉ：０．２ 〜２．０ ％
Ｓｉは、固溶強化成分として、強度−伸びバランスを改善しつつ強度の向上に有効に寄与するが、過剰な添加は、溶融亜鉛めっきの合金化を阻害し、延性や溶接性を劣化させるので、Ｓｉは ２．０％以下で含有させるものとした。また、Ｓｉは、フェライト相の安定化およびオーステナイトへの炭素の濃化促進効果があるため、その含有量の下限を０．２ ％とした。
【００２２】
Ｍｎ：１．０ 〜３．０ ％および／またはＮｉ：０．５ 〜３．０ ％
ＭｎおよびＮｉはいずれも、オーステナイト安定化元素であり、Ａ_１ ，Ａ_３ 変態点を低下させる作用を通じて結晶粒の微細化に寄与し、また第２相の形成を進展させる作用を通じて強度−延性バランスを高める作用を有する。しかしながら、Ｍｎ含有量が １．０％に満たない、またはＮｉ含有量が ０．５％に満たないと、安定したオーステナイト相の形成と所望の強度が得られず、一方多量の添加は鋼の硬質化を招き、却って強度−延性バランスを劣化させるので、Ｍｎ， Ｎｉ含有量はいずれも ３．０％以下で含有させるものとした。
【００２３】
Ｔｉ：０．２ ％以下および／またはＮｂ：０．２ ％以下
Ｔｉ， Ｎｂを添加することによって、ＴｉＣやＮｂＣ等が析出し、鋼板の再結晶温度が上昇する効果がある。これらは各々単独で添加しても複合して添加してもよいが、いずれも ０．２％を超えて添加しても効果が飽和するだけでなく、析出物が多くなりすぎてフェライトの延性の低下を招くので、いずれも ０．２％以下で含有させるものとした。
【００２４】
Ａｌ：０．０１〜０．１ ％
Ａｌは、脱酸剤として作用し、鋼の清浄度に有効な元素であり、脱酸の工程で添加することが望ましい。ここに、Ａｌ量が０．０１％に満たないとその添加効果に乏しく、一方 ０．１％を超えると効果は飽和し、むしろ製造コストの上昇を招くので、Ａｌは０．０１〜０．１ ％の範囲に限定した。
【００２５】
Ｐ：０．１ ％以下
Ｐは、延性の大きな低下を招くことなく安価に高強度化を達成する上で有用な元素であるが、含有量が ０．１％を超えると加工性や靱性の低下を招くので、Ｐの含有量は ０．１％以下に限定した。なお、加工性や靱性に対する要求が厳しい場合には、Ｐはむしろ低減させることが好ましいので、この場合には０．０２％以下とすることが望ましい。
【００２６】
Ｓ：０．０２％以下
Ｓは、熱延時における熱間割れの原因になるだけでなく、鋼板中にＭｎＳ等の介在物として存在し延性や穴拡げ加工性の劣化を招くので、極力低減することが望ましいが、０．０２％までは許容できるので、本発明では０．０２％以下とした。より好ましくは ０．００５％以下である。
【００２７】
Ｎ：０．００５ ％以下
窒素は、時効劣化をもたらす他、降伏延びの発生を招くことから、０．００５ ％以下に抑制するものとした。
【００２８】
以上、基本成分について説明したが、本発明ではその他にも、以下に述べる元素を適宜含有させることができる。
Ｍｏ：１．０ ％以下およびＣｒ：１．０ ％以下のうちから選んだ一種または二種
Ｍｏ，Ｃｒはいずれも、強化成分として、必要に応じて含有させることができるが、多量の添加はかえって強度−延性バランスを劣化させるので、それぞれ １．０％以下で含有させることが望ましい。なお、上記の作用を十分に発揮させるには、Ｍｏ， Ｃｒはそれぞれ０．０１％以上含有させることが好ましい。
【００２９】
Ｃａ， ＲＥＭ およびＢのうちから選んだ一種または二種以上を合計で ０．００５％以下
Ｃａ， ＲＥＭ，Ｂはいずれも、硫化物の形態制御や粒界強度の上昇を通じて加工性および伸びフランジ性を改善する効果を有しており、必要に応じて含有させることができる。しかしながら、過剰な含有は清浄度に悪影響を及ぼすおそれがあるため、合計で ０．００５％以下とするのが望ましい。なお、上記した作用を十分に発揮させるためには、 Ｃａ， ＲＥＭ， Ｂのうちから選んだいずれか一種または二種以上を合計で０．０００５％以上含有させることが好ましい。
【００３０】
以上、適正な成分組成範囲について説明したが、本発明では各成分が上記の組成範囲を単に満足しているだけでは不十分で、Ｃ，Ｓｉ， Ｍｎ， Ｎｉ， ＴｉおよびＮｂについては、下記（１）， （２）， （３） 式をそれぞれ満足する範囲で含有させる必要がある。

【００３１】
なお、上記のＡ_１ ， Ａ_３ はそれぞれ、鋼のＡｃ_１変態点温度（℃）、Ａｃ_３変態点温度（℃）の予測値であり、発明者らの詳細な基礎実験により導出された成分回帰式である。この予測値温度（℃）は、２℃／ｓ以上、２０℃／ｓ以下の昇温速度で加熱する際に適用して特に好適である。
【００３２】
以下、上記の（１）， （２）， （３） 式の限定理由を順に説明する。
（１） 式は、Ｔｉ，Ｎｂの添加量を規定する条件であり、以下の知見に基づく。
一般に、Ｔｉ，Ｎｂを添加するとＴｉＣやＮｂＣ等が析出し、鋼板の再結晶温度が上昇する効果があることが知られている。そこで、Ｔｉ，Ｎｂ添加量と再結晶温度Ｔｒｅの関係について詳細に調査したところ、Ｔｉ，Ｎｂをある量以上添加すると、再結晶温度は上記（６） 式で算出されるＡ_３ と等価になることが判明した。
【００３３】
図１に、Ａ_１ ＝７００ ℃、Ａ_３ ＝８５５ ℃に調整した鋼組成において、Ｔｉ，Ｎｂ添加量を種々に変更した場合のＴｉ，Ｎｂ添加量と再結晶温度Ｔｒｅとの関係について調べた結果を示す。なお、ここで再結晶温度Ｔｒｅは、加熱温度を種々に変化させて連続焼鈍を実験室的に行い、硬度を測定すると共に組織を観察することにより決定した。また、Ｔｉ添加量はＴｉＣを析出させる上での有効Ｔｉ量としてＴｉ^＊ を用い、Ｎｂ添加量はＴｉに換算するため ４８／９３・［％Ｎｂ］ を用いて、Ｔｉ， Ｎｂ添加量と再結晶温度との関係について表わしている。
同図によれば、 ６３７．５＋４９３０｛Ｔｉ^＊ ＋ （４８／９３）・［％Ｎｂ］ ｝が ７００℃すなわちＡ_１ 以上になると、再結晶温度Ｔｒｅは ８５５℃近傍すなわちＡ_３ 近傍に急上昇し飽和することが分かる。
【００３４】
次に、図２に、 ６３７．５＋４９３０｛Ｔｉ^＊ ＋ （４８／９３）・［％Ｎｂ］ ｝≧Ａ_１ の条件下において、Ａ_３ （Ｃ，Ｓｉ，Ｍｎ， Ｎｉ等を変化させることで変動）を種々に変化させた場合におけるＡ_３ と再結晶温度Ｔｒｅとの関係について調べた結果を示す。
同図に示したとおり、 ６３７．５＋４９３０｛Ｔｉ^＊ ＋ （４８／９３）・［％Ｎｂ］ ｝≧Ａ_１ の条件下では、再結晶温度ＴｒｅはＡ_３ と等価になっている。
【００３５】
この理由については、必ずしも明確ではないが、以下のように考えられる。
すなわち、Ｔｉ，Ｎｂが添加され、それらの微細炭化物のピン止め力により再結晶温度が上昇し、Ａ_１ 未満のフェライト（α）域で再結晶できなくなった場合、未再結晶の加工αのまま（フェライト＋オーステナイト（γ））２相域温度になり、高転位密度部、不均一変形部などの優先核生成サイトにおいて、加工αからの再結晶α核生成とα→γ変態核生成の競合が生じる。この時、α→γ変態の駆動力の方が再結晶の駆動力よりも大きいため、再結結晶α核生成より優先してγ核が次々と生成し、優先核生成サイトを占有すると考えられる。
このα→γ変態での原子再配列により歪み（転位）は消費され、転位密度の低い加工αのみ残留し、加工αの再結晶はますます困難となる。温度が上昇し、Ａ_３ を超え、γ単相域になって初めて歪みが完全に解消され、見かけ上再結晶が完了する。これが、再結晶温度がＡ_３ に一致し、飽和する機構と考えられる。
なお、この際のα→γ変態は、加工α（優先核生成サイトが多い）から核生成することになるので、再結晶が完了した高温でのγ粒は微細化する。従って、焼鈍中の高温γ粒微細化のために再結晶温度をＡ_３ とすることは有効であるので、本発明では式（１） を満足するＴｉ， Ｎｂを添加することにしたのである。
【００３６】
次に、 （２）式は、Ａ_３ を規定する条件である。
上述したとおり、 （１）式を満足する場合には、Ａ_３ は実質的に再結晶温度になるため、Ａ_３ 以上の温度で再結晶焼鈍を行う必要がある。ここに、Ａ_３ が ８６０℃を超えた場合、再結晶焼鈍温度をより高温で施す必要が生じ、γ粒成長が激しく、結果として平均結晶粒径：３．５ μｍ 以下の微細粒は得られなかった。よって、Ａ_３ ≦８６０ ℃を満足させる必要がある。
【００３７】
次に、 （３）式は、ＭｎやＮｉすなわちオーステナイト安定化元素の添加量を規定する条件である。
オーステナイト安定化元素の増大により、ＣＣＴ 図におけるフェライトスタート線が低温側にシフトすることにより、焼鈍後の冷却過程におけるγ→α変態時の変態過冷度が増大してαが微細核生成することにより、α結晶粒が微細化する。ここに、平均結晶粒径：３．５ μｍ 以下の微細粒を得るためには、上掲した（１）， （２）式に加えて ［％Ｍｎ］＋［％Ｎｉ］ ≧ １．３（％）とする必要があった。
なお、 ［％Ｍｎ］＋［％Ｎｉ］ ≧ １．３（％）さえ満足していれば、ＭｎやＮｉは単独添加でも複合添加でもどちらでも良い。より好ましくは ［％Ｍｎ］＋［％Ｎｉ］ ≧ ２．０（％）の範囲である。
【００３８】
次に、鋼組織について説明する。
本発明において、鋼組織は、フェライト相の体積分率を７０ ｖｏｌ％以上、残留オーステナイトの体積分率を５ ｖｏｌ％以上にすると共に、フェライトの平均結晶粒径を ３．５μｍ 以下とする必要がある。
すなわち、フェライトは軟質な組織であり、延性や加工性の向上に寄与するだけでなく、炭素を固溶し難いため、めっき前の熱処理時にオーステナイトに炭素を濃化させるのに有効に寄与する。従って、本発明で所期した強度、延性を有する溶融亜鉛めっき鋼板を得るためには、微細フェライトを主体とする鋼組織とする必要があり、特に平均結晶粒径が ３．５μｍ 以下の微細フェライト相の体積分率を７０ ｖｏｌ％以上とすることが重要だからである。
ここに、フェライトの平均結晶粒径が ３．５μｍ を超えると強度−伸びバランスが低下すると共に、伸びフランジ性が劣化する。また、フェライトの体積分率が７０ ｖｏｌ％に満たないと、延性が著しく低下し、加工性が劣化する。
【００３９】
また、第２相は、残留オーステナイト相とする。この残留オーステナイト相は、鋼板の延性に大きく寄与する組織であり、その効果を得るためには、体積分率で組織全体の５ ｖｏｌ％以上とする必要がある。
【００４０】
なお、上記したフェライト相および残留オーステナイト相の他に、ベイナイト相やパーライト相、さらにはマルテンサイト相が生成する場合があるが、これらの相の合計が体積分率で２０ ｖｏｌ％以下であれば、許容できる。
【００４１】
次に、製造条件について説明する。
上記の好適成分組成に調整した鋼を、転炉などで溶製し、連続鋳造法等でスラブとする。この鋼素材であるスラブを、高温状態のまま、あるいは冷却したのち、１２００℃以上に加熱してから、熱間圧延を施し、ついで冷間圧延後、温度Ａ_３ （℃）以上、（Ａ_３ ＋３０）（℃）以下で再結晶焼鈍を施し、その後好ましくは ３００℃までを１０℃／ｓ以上の速度で冷却する。
【００４２】
上記の工程において、スラブの加熱温度が１２００℃未満では、ＴｉＣなどが十分に固溶せずに粗大化し、後の再結晶焼鈍工程での再結晶温度上昇効果および結晶粒成長抑止効果が不十分となるため、スラブの加熱温度は１２００℃以上とする必要がある。
また、本発明において、上記スラブ加熱温度以外の熱間圧延条件は特に制限されるものではなく、常法に従えばよい。なお、熱間仕上げ圧延出側温度は特に限定されるものではないが、Ａｒ_３変態点未満では、圧延中にαとγが生じて、鋼板にバンド状組織が生成し易くなり、かかるバンド状組織は冷間圧延後や焼鈍後にも残留し、材料特性に異方性を生じさせる原因となる場合があるので、熱間仕上げ圧延終了温度はＡｒ_３変態点以上とすることが好ましい。
【００４３】
熱間圧延終了後の巻取り温度も特に限定されるものではないが、５００ ℃未満または ６５０℃超えでは、窒素による時効劣化を抑制するためのＡｌＮの析出が不十分であり、材料特性が劣ることとなる。また、鋼板の組織を均一化し、その結晶粒径をなるべく微細で均一化するためにも、コイルの巻取り温度は ５００℃以上、６５０ ℃以下とすることが好ましい。
【００４４】
ついで、好ましくは熱延鋼板表面の酸化スケールを酸洗により除去したのち、冷間圧延に供して、所定の板厚の冷延鋼板とする。ここに、酸洗条件や冷間圧延条件は特に制限されるものでなく、常法に従えばよい。
なお、冷間圧延時の圧下率は、再結晶焼鈍時の核生成サイトを増やし、結晶粒の微細化を促すという観点から４０％以上とすることが望ましく、一方圧下率を上げすぎると鋼板の加工硬化によって操業が困難となるので、圧下率の上限は９０％以下程度とするのが好ましい。
【００４５】
ついで、得られた冷延鋼板を、前掲（６） 式に示した温度Ａ_３（℃）以上、（Ａ_３＋３０）（℃）以下に加熱して、再結晶焼鈍を施す。
前述のように成分調整した本発明の鋼素材では、Ａ_３ が実質的に再結晶温度と等価となっているので、Ａ_３ 未満の温度では再結晶が不十分となる。一方、（Ａ_３ ＋３０）（℃）を超える温度では、焼鈍中のγ粒の成長が激しく、微細化に不適切である。この再結晶焼鈍は、連続焼鈍ラインで行うことが好ましく、連続焼鈍する場合の焼鈍時間は再結晶が生じる１０秒から １２０秒程度とすることが好ましい。というのは、１０秒より短時間では再結晶が不十分であり、圧延方向に伸展したままの加工組織、再結晶していない回復組織が残存するために、十分な延性が確保できない場合があり、一方 １２０秒より長時間ではγ結晶粒の粗大化を招いて、所望の強度を得ることができないことがあるからである。
【００４６】
なお、再結晶焼鈍時の昇温速度は通常、連続焼鈍の行われる２〜２０℃／ｓ程度とすればよく、また再結晶焼鈍後の冷却は、結晶粒の成長を抑制するため、焼鈍温度から ３００℃までの平均冷却速度を１０℃／ｓ以上として冷却することが好ましい。
【００４７】
ついで、上記の再結晶焼鈍後、酸洗を行い、めっき性に悪影響を及ぼす表面酸化物を除去する。すなわち、焼鈍中に鋼板表面に析出したＰ，Ｓｉ， Ｍｎ， Ｃｒなどが酸化物として濃化した表面濃化層を除去する。なお，このような除去すべき表面濃化層は、軽酸洗によって除去可能であるので、従来から行われている連続溶融亜鉛めっき処理前の軽酸洗で十分である。
【００４８】
つぎに、Ａ_１ 以上、（Ａ_１ ＋７０）℃以下の温度域にて５〜３０秒間加熱した後、溶融亜鉛めっき開始温度まで好ましくは５〜１５℃／ｓの速度で冷却する。このような条件で熱処理することにより、Ｃ，Ｍｎ，Ｍｏ，Ｎｉなどの置換型の合金元素が変態により生成するオーステナイト相へと濃化して安定化し、その結果、冷却して得られる残留オーステナイト相を室温でも安定化することができ、良好な延性を得ることができる。
【００４９】
本発明では、前述したように、上記熱処理に先行して行う再結晶焼鈍で鋼の結晶粒を微細化しているため、Ｃや他の合金元素がγ相へと移動する距離が短く、オーステナイトへのＣなどの合金元素が濃化し易いので、残留オーステナイトを安定して得ることができる。
この熱処理における加熱温度がＡ_１ 未満では、オーステナイト相が生成しないため、冷却後に残留オーステナイト相が安定して得られない。一方、加熱温度が（Ａ_１ ＋７０）℃超では、結晶粒成長が進行し、冷却後の微細組織が得られなくなる。よって、加熱温度はＡ_１ 以上、（Ａ_１ ＋７０） ℃以下とした。また、加熱時間が５秒未満では変態に要する時間が短かすぎてオーステナイトの生成に不利であり、一方３０秒より長時間ではオーステナイト中の炭素量が平衡量に達し、その効果がなくかえってオーステナイト中の炭素量が減少してしまい、安定な残留オーステナイトを生成する上で不利となる。従って、熱処理の時間は５〜３０秒に限定した。
【００５０】
また、熱処理時後の冷却速度は、結晶粒成長を抑制するためには５℃／ｓ以上とすることが好ましい。しかしながら、冷却速度があまりに速すぎるとマルテンサイトが生成して延性が低下するおそれがあるので、１５℃／ｓ以下とすることが好ましい。
【００５１】
なお、上記の熱処理は連続溶融亜鉛めっきラインにて行うことが望ましい。
上記熱処理後の冷却に引き続き溶融亜鉛めっきを行い、あるいはさらに合金化処理を施して、フェライトを７０ ｖｏｌ％以上、残留オーステナイトを５ ｖｏｌ％以上有する鋼組織とし、かつフェライト粒径が３．５ μｍ 以下の高張力溶融亜鉛めっき冷延鋼板を得る。
【００５２】
また、溶融亜鉛めっき処理あるいは合金化処理は通常の条件で行えばよい。
めっき温度が不必要に高温になりすぎると、残留オーステナイト量の確保に不利となり、一方温度が低いと亜鉛浴中での亜鉛と鉄の反応が遅くなって生産性が低下するため、溶融亜鉛めっき時の板温は ４５０〜５５０ ℃程度とするのが有利である。
【００５３】
さらに、溶融亜鉛めっき後の合金化処理に際し、あまりに高温では残留オーステナイトの確保に不利となり、一方低温すぎると合金化の進行が遅くなり、生産性の低下を招くので、合金化処理時の板温は ５００〜５６０ ℃程度とするのが望ましい。
【００５４】
【実施例】
表１に示す成分組成になるスラブを、表２に示す条件でスラブ加熱後、常法に従い熱間圧延して ４．０ｍｍ厚の熱延板とした。この時、熱間仕上げ圧延出側温度はＡｒ_３変態点以上とし、巻き取り温度は ６００℃とした。この熱延板を、酸洗後、冷間圧延（圧下率：６０％）して、１．６ ｍｍ厚の冷延板としたのち、連続焼鈍ラインにて同じく表２に示す条件下で焼鈍を行い、ついで酸洗後、連続溶融亜鉛めっきラインにて、熱処理に引き続きめっき処理さらには合金化処理を施して、溶融亜鉛めっき鋼板とした。冷延板焼鈍時の昇温速度は５〜１０℃／ｓ、冷却速度は ３００℃までの平均冷却速度を２０〜２５℃／ｓとした。連続溶融亜鉛めっきラインでのめっき前加熱温度から溶融亜鉛めっき温度までの冷却速度は５〜１５℃／ｓであった。また、溶融亜鉛めっき処理での浴中の鋼板温度は４６５ ℃、その後合金化処理を施した場合の合金化温度は ５００℃であった。なお、Ｎｏ．１６ は、溶融亜鉛めっきのみとし、合金化処理は施さなかった。
かくして得られた溶融亜鉛めっき鋼板の組織、引張特性について調べた結果を表３に示す。
【００５５】
なお、組織は、鋼板の圧延方向断面について、光学顕微鏡あるいは電子顕微鏡を用いて調べ、併せてフェライトの体積率および平均結晶粒径を測定した。ここで、フェライトの平均結晶粒径はＪＩＳ Ｇ ０５５２に規定される切断法に準拠して求めた。
また、引張特性（引張強さ（ＴＳ）、伸び（ＥＬ））は、鋼板の圧延方向から採取したＪＩＳ５号試験片を用いた引張試験により測定した。
残留オーステナイト量は、鋼板を板厚方向に板厚の１／４ 位置まで研磨し、この１／４ 面でのＸ線回折強度を測定して求めた。入射Ｘ線にはＭｏＫαを使用し、フェライト相の｛１１０｝、｛２００｝、｛２１１｝各面のＸ線回折強度に対する残留オーステナイト相の｛２００｝、｛２２０｝、｛３１１｝各面のＸ線回折強度を求め、これらの平均値を残留オーステナイトの体積率とした。
延びフランジ性は、以下に述べる穴拡げ試験によって評価した。すなわち、日本鉄鋼連名規格ＪＦＳＴ１００１に準じて採取した試験片に、１０ｍｍφ（Ｄ_０ ）の打ち抜き穴を加工したのち、頂角：６０°の円錐ポンチで押し拡げる加工を施し、割れが板厚を貫通した直後の穴径Ｄ（ｍｍ）を求め、次式
λ＝｛（Ｄ−Ｄ_０）／Ｄ_０）｝×１００ ％
で得られる穴拡げ率λで評価した。
【００５６】
【表１】

【００５７】
【表２】

【００５８】
【表３】

【００５９】
表３に示したとおり、発明例はいずれも、フェライトの体積分率が７０ ｖｏｌ％以上、残留オーステナイトの体積分率が５ ｖｏｌ％以上で、かつフェライトの平均結晶粒径が ３．５μｍ 以下と微細であり、またいずれもＴＳ×ＥＬが ２００００ ＭＰａ・％以上、ＴＳ×λが ６７０００ ＭＰａ・％以上と強度−延性バランスおよび強度−穴拡げバランスに優れていることが分かる。
これに対し、Ｎｏ．４は、めっき前加熱温度が適正上限温度（７３２ ℃）を大きく超えたため、結晶粒成長が激しくて粗大化したため、ＴＳ×ＥＬが低下した。
Ｎｏ．９は、スラブの加熱温度が低かったため、ＴｉＣが粗大化し、再結晶温度上昇効果が抑制されて鋼板の結晶粒微細化効果が得られず、結晶粒径が大きくなった。また、ＴＳ×ＥＬ、ＴＳ×λ値も小さくなっている。
Ｎｏ．１０ は、焼鈍温度が本発明の適正上限温度（８８４ ℃）を大きく超えたため、結晶粒成長が激しく、ＴＳ×ＥＬ値が劣化した。
Ｎｏ．１１ は、焼鈍温度が本発明の下限（８５４ ℃）に満たなかったため、再結晶が完了せず、加工組織が残留したため、ＴＳ×ＥＬ値が劣っている。
Ｎｏ．１２ は、めっき前の加熱温度がＡ_１ 未満であったため、焼き戻されただけの組織となり、オースナイト相が形成されず、延性の低下を招いた。
Ｎｏ．２１ は、（１） 式を満足するにはＴｉ， Ｎｂ量が不足しているため、再結晶が低温（α域）で起こり、再結晶温度とＡ_３ が一致しない。このため、冷却後の焼鈍時に再結晶が通常のようにα域で生じ、結晶粒が粗大化して、十分な強度が得られず、伸び、穴広げ率も低い値であった。
Ｎｏ．２２ は、Ａ_３ が ８６０℃以上であることから、高温焼鈍が必要となり、その結果結晶粒が成長して、ＴＳ×ＥＬ値が劣化した。
Ｎｏ．２３ は、（Ｎｉ＋Ｍｎ）量が少ないために、焼鈍後冷却過程でのγ−α変態時の過冷度が小さく、フェライトが微細核生成することができなかったため、結晶粒が粗大化した。
【００６０】
【発明の効果】
かくして、本発明によれば、超微細粒組織を有し、延性に優れた高張力冷延鋼板を、製造設備の大幅な改造を伴うことなしに安定して製造することができ、産業上極めて有用である。
【図面の簡単な説明】
【図１】Ａ_１ ＝７００ ℃、Ａ_３ ＝８５５ ℃に調整した鋼組成において、Ｔｉ，Ｎｂ添加量を種々に変更した場合のＴｉ，Ｎｂ添加量と再結晶温度との関係を示した図である。
【図２】６３７．５＋４９３０｛Ｔｉ^＊ ＋ （４８／９３）・［％Ｎｂ］ ｝≧Ａ_１ の条件下において、Ａ_３ を種々に変化させた場合におけるＡ_３ と再結晶温度Ｔｒｅとの関係を示した図である。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a high-strength hot-dip galvanized steel sheet and a method for producing the same, which are suitable for use as a steel for automobiles and home appliances, and also for machine structural use. It is intended to achieve an advantageous improvement.
[0002]
[Prior art]
Steel materials used for automobiles, home appliances, and steel sheets for machine structures are required to have excellent mechanical properties such as strength and workability and excellent corrosion resistance.
In this regard, hot-dip galvanized steel sheet has not only high corrosion resistance, but also has the advantage that it can be manufactured at extremely low cost by a continuous hot-dip galvanizing line that can perform recrystallization annealing and galvanizing processing on the same line. I have.
[0003]
In recent years, with respect to high-strength steels, the target is shifting to the development of high-strength steel sheets that can achieve both high performance characteristics and low cost.
Further, steel sheets for automobiles are required to have not only high strength but also excellent impact resistance in order to protect occupants in the event of a collision.
[0004]
In addition, since most automotive parts made of steel sheets are formed by press working, excellent formability is required for steel parts for automotive parts. In particular, in the case of components constituting members such as a skeleton member for ensuring the strength of an automobile body, reinforcement, and the like, component molding utilizing ductility and stretch flange deformation is often performed.
However, in general, strength and ductility are contradictory properties, and the ductility decreases as the strength increases. For this reason, steel sheets for automobile parts are strongly required to have high ductility and good ductility at the same time.
[0005]
In order to satisfy these demands, a method has been proposed in which hot-dip galvanizing is performed while retaining retained austenite, which contributes to the ductility of a steel sheet, using work-induced transformation (for example, see Patent Document 1).
In this method, a steel sheet adjusted to a predetermined component composition is supplied to a continuous hot-dip galvanizing line (Ac ₁ Above + 30 ° C) ₃ Annealing for 30 seconds or more in the temperature range below the temperature ₁ (Point + 20 ° C) ~ Ac ₁ It is cooled at a cooling rate of 5 ° C./s or less to a point of 5 ° C., and then cooled to a temperature of 520 ° C. or less at a cooling rate of 6 ° C./s or less. In this method, the steel is allowed to stay in a temperature range of 520 to 400 ° C. for 90 seconds to 300 seconds, and then cooled to 200 ° C. or less to contain 3% or more by volume of retained austenite.
[0006]
However, in this method, in order to concentrate carbon in austenite existing at the time of annealing and to further stabilize the carbon, (Ac ₁ (Point + 20 ° C) ~ Ac ₁ There is a problem that the steel sheet has to be cooled within a very narrow temperature range of points, and it is extremely difficult to control the temperature of the steel sheet.
Also, (Ac ₁ Above + 30 ° C) ₃ From the temperature range below the point, (Ac ₁ (Point + 20 ° C) ~ Ac ₁ The aim of cooling to the temperature range of the point is to promote fermentation by growing ferrite crystal grains in order to concentrate carbon in austenite. However, ductility and stretch flangeability (also referred to as hole expanding properties) It is not desirable to grow the crystal grains from the viewpoint of improving the grain size and obtaining a high-strength steel sheet.
[0007]
In addition, there has been proposed a method in which hot-dip galvanizing is performed while retaining retained austenite, which contributes to ductility of a steel sheet, by utilizing work-induced transformation (for example, see Patent Document 2).
According to this method, a steel sheet adjusted to a predetermined component composition is obtained by (Ac ₃ (-80 ° C. point) After performing a primary heat treatment for 5 seconds or more in the above temperature range, a primary step of cooling to a temperature of Ms point or less at a cooling rate of 5 ° C./s or more and less than 10 ° C./s, (Ac ₁ Transformation point ~ Ac ₃ (Transformation point) after performing a secondary heat treatment for 5 seconds or more in the temperature range, cooling at a cooling rate of 5 ° C./s or more to a temperature range of 500 ° C. or less, and then performing a hot-dip galvanizing process. After forming a hot-dip galvanized layer on the surface of the steel sheet, a tertiary step of cooling to 300 ° C. at a cooling rate of 5 ° C./s or more, and further reheating to a temperature range of 450 to 550 ° C. By cooling to 300 ° C. at a cooling rate of 5 ° C./s or more, 2% or more by volume of retained austenite is to be contained.
However, this method is not only extremely difficult to control the heating temperature, the cooling rate and the cooling stop temperature for defining the content of ferrite, martensite and residual austenite in the steel sheet, but also includes massive martensite, Furthermore, it is difficult to improve both ductility and hole-expandability, in part because the average crystal grain size needs to be 5 μm or more.
[0008]
On the other hand, as a method for obtaining a fine grain structure, a large rolling reduction method has been conventionally known. The key point of the mechanism for refining the structure in this large rolling reduction method is to apply a large reduction to austenite grains to promote γ-α strain-induced transformation (for example, see Patent Document 3).
Also, a case where a controlled rolling method or a controlled cooling method is applied is known (for example, see Patent Document 4).
[0009]
In addition, regarding the material steel, a steel structure at least partially composed of ferrite is formed, and the transformation point (Ac ₁ Point) The temperature is raised to the above temperature range, or Ac ₁ After a part of or all of the structure is once transformed back to austenite by holding it in the temperature range above the temperature for a certain period of time, ultra-fine austenite grains appear, and then cooled to obtain an isotropic crystal having an average crystal grain size of 5 μm or less. A technique of forming a structure mainly composed of crystalline ferrite grains has been proposed (Patent Document 5).
[0010]
All of the techniques described above are techniques for refining crystal grains in the hot rolling process, that is, techniques for refining hot rolled sheets.
In this regard, compared to hot-rolled steel sheets, the thickness is thin, and the requirements for thickness accuracy and surface properties are strict. In addition, for hot-dip galvanized steel sheets to be subjected to surface treatment, ordinary cold rolling-annealing-hot-dip zinc Few techniques for refining crystal grains in the plating process are found.
[0011]
[Patent Document 1]
JP-A-11-131145 (Claims)
[Patent Document 2]
JP 2001-192767 A (Claims)
[Patent Document 3]
Japanese Patent Publication No. 5-65564 (Claims)
[Patent Document 4]
JP-A-63-128117 (Claims)
[Patent Document 5]
JP-A-2-301540 (Claims)
[0012]
[Problems to be solved by the invention]
The present invention has been developed in view of the above-mentioned current situation, and improves strength and ductility of a surface-treated steel sheet used as a steel sheet for automobiles, home appliances and mechanical structures by enabling ultra-fine graining. It is an object of the present invention to propose a high-strength hot-dip galvanized steel sheet having good hot-dip galvanizing properties, together with its advantageous production method.
The hot-dip galvanized steel sheet referred to in the present invention includes a so-called non-alloyed hot-dip galvanized steel sheet (GI steel sheet) that is not subjected to alloying treatment after hot-dip galvanizing, and a so-called hot-dip galvanized steel sheet that is subjected to alloying treatment after hot-dip galvanizing. It means both plated steel sheets (GA steel sheets).
[0013]
[Means for Solving the Problems]
By the way, the present inventors have conducted intensive studies to solve the above-described problems with respect to a high-strength hot-dip galvanized steel sheet using a cold-rolled material as a base plate, that is, a high-strength hot-dip galvanized cold-rolled steel sheet. Adjust the recrystallization temperature and A ₁ And A ₃ By controlling the transformation temperature and optimizing the recrystallization annealing temperature after cold rolling and the subsequent cooling rate, an ultrafine grain structure having an average crystal grain size of 3.5 μm or less can be obtained. ₁ Above the transformation point, Ac ₃ It has been found that heat treatment in a temperature range below the transformation point significantly improves ductility.
The present invention is based on the above findings.
[0014]
That is, the gist configuration of the present invention is as follows.
1. In mass%,
C: 0.03 to 0.16%,
Si: 0.2 to 2.0%,
Mn: 1.0 to 3.0% and / or Ni: 0.5 to 3.0%,
Ti: 0.2% or less and / or Nb: 0.2% or less;
Al: 0.01 to 0.1%,
P: 0.1% or less,
S: 0.02% or less and
N: 0.005% or less
And C, Si, Mn, Ni, Ti and Nb are contained in ranges satisfying the following formulas (1), (2) and (3), and the balance is Fe and unavoidable impurities. Has a steel structure with a volume fraction of 70 vol% or more, a retained austenite volume fraction of 5 vol% or more, and an average crystal grain size of the ferrite of 3.5 μm or less, and further has a hot-dip galvanized surface. A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility characterized by having a layer.

[0015]
2. In the above item 1, the steel sheet may further comprise
Mo: 1.0% or less and
Cr: 1.0% or less
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that it has a composition containing one or two selected from the above.
[0016]
3. In the above 1 or 2, the steel sheet further comprises, by mass%,
One or more kinds selected from Ca, REM and B are 0.005% or less in total
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized by having a composition containing:
[0017]
4. In mass%,
C: 0.03 to 0.16%,
Si: 0.2 to 2.0%,
Mn: 1.0 to 3.0% and / or Ni: 0.5 to 3.0%,
Ti: 0.2% or less and / or Nb: 0.2% or less;
Al: 0.01 to 0.1%,
P: 0.1% or less,
S: 0.02% or less and
N: 0.005% or less
And a steel material containing C, Si, Mn, Ni, Ti and Nb in ranges satisfying the following formulas (1), (2) and (3), and the balance being Fe and unavoidable impurities. Is heated to 1200 ° C. or higher, hot-rolled, then cold-rolled, and then subjected to a temperature A determined by the following equation (6). ₃ ℃ or more, (A ₃ +30) ° C. or lower, followed by pickling, followed by A ₁ ℃ or more, (A ₁ +70) a hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, which is subjected to a heat treatment for 5 to 30 seconds in a temperature range of not more than ° C and subsequently to a hot-dip galvanizing treatment or an alloying treatment. Manufacturing method.

[0018]
5. In the above item 4, the steel material further contains, by mass%,
Mo: 1.0% or less and
Cr: 1.0% or less
A method for producing a hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that the composition contains one or two selected from the group consisting of:
[0019]
6. In the above 4 or 5, the steel material further contains, by mass%,
One or more kinds selected from Ca, REM and B are 0.005% or less in total
A method for producing a hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized by having a composition containing
[0020]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described specifically.
First, the reason why the composition of the steel material is limited to the above range in the present invention will be described. In addition, "%" display about a component shall mean the mass% unless there is particular notice.
C: 0.03 to 0.16%
C is not only an inexpensive strengthening component, but also a useful element for stabilizing retained austenite that contributes to ductility. However, if the content is less than 0.03%, the effect of the addition is poor. On the other hand, if the content is more than 0.16%, the weldability is deteriorated, so that C is in the range of 0.03 to 0.16%. Limited.
[0021]
Si: 0.2 to 2.0%
Si, as a solid solution strengthening component, effectively contributes to the improvement of strength while improving the strength-elongation balance, but excessive addition inhibits alloying of hot-dip galvanizing and deteriorates ductility and weldability. , And Si are contained at 2.0% or less. Since Si has the effect of stabilizing the ferrite phase and accelerating the concentration of carbon into austenite, the lower limit of the content is set to 0.2%.
[0022]
Mn: 1.0 to 3.0% and / or Ni: 0.5 to 3.0%
Mn and Ni are both austenite stabilizing elements, ₁ , A ₃ It has the effect of contributing to the refinement of crystal grains through the action of lowering the transformation point and of increasing the strength-ductility balance through the action of promoting the formation of the second phase. However, if the Mn content is less than 1.0% or the Ni content is less than 0.5%, a stable austenite phase cannot be formed and the desired strength cannot be obtained. Since Mn and Ni are hardened and the strength-ductility balance is rather deteriorated, the content of Mn and Ni is set to 3.0% or less.
[0023]
Ti: 0.2% or less and / or Nb: 0.2% or less
By adding Ti and Nb, TiC, NbC and the like are precipitated, and there is an effect that the recrystallization temperature of the steel sheet increases. Each of these may be added alone or in combination. However, adding more than 0.2% not only saturates the effect, but also increases the amount of precipitates and the ductility of ferrite. In each case, the content was 0.2% or less.
[0024]
Al: 0.01 to 0.1%
Al acts as a deoxidizing agent and is an element effective for the cleanliness of steel, and is desirably added in the deoxidizing step. Here, if the Al content is less than 0.01%, the effect of the addition is poor, while if it exceeds 0.1%, the effect is saturated and rather increases the production cost. Limited to the 1% range.
[0025]
P: 0.1% or less
P is an element useful for achieving high strength at low cost without causing a large decrease in ductility. However, if the content exceeds 0.1%, workability and toughness are reduced. The content was limited to 0.1% or less. In the case where demands on workability and toughness are severe, it is preferable to reduce P, and in this case, it is preferable to set the content to 0.02% or less.
[0026]
S: 0.02% or less
S is not only a cause of hot tearing at the time of hot rolling, but also exists as an inclusion such as MnS in a steel sheet and causes deterioration of ductility and hole expanding workability. Therefore, it is desirable to reduce S as much as possible. Since up to 02% is acceptable, it is set to 0.02% or less in the present invention. More preferably, it is 0.005% or less.
[0027]
N: 0.005% or less
Nitrogen not only causes aging deterioration but also causes yield elongation, so it was limited to 0.005% or less.
[0028]
As described above, the basic components have been described. However, in the present invention, other elements described below can be appropriately contained.
One or two selected from Mo: 1.0% or less and Cr: 1.0% or less
Both Mo and Cr can be contained as necessary as a strengthening component. However, since the addition of a large amount rather deteriorates the strength-ductility balance, it is desirable that each of Mo and Cr be contained at 1.0% or less. In order to sufficiently exhibit the above-mentioned effects, it is preferable that Mo and Cr are each contained in an amount of 0.01% or more.
[0029]
One or more selected from Ca, REM and B are 0.005% or less in total
Ca, REM, and B all have the effect of improving workability and stretch flangeability through morphological control of sulfides and increase in grain boundary strength, and can be contained as necessary. However, an excessive content may adversely affect the cleanliness. Therefore, the total content is desirably 0.005% or less. In order to sufficiently exert the above-mentioned effects, it is preferable to contain one or more selected from Ca, REM, and B in a total amount of 0.0005% or more.
[0030]
Although the appropriate component composition range has been described above, in the present invention, it is not sufficient that each component simply satisfies the above composition range. For C, Si, Mn, Ni, Ti and Nb, the following ( It is necessary to include the formulas (1), (2), and (3) in a range satisfying each.

[0031]
Note that the above A ₁ , A ₃ Is the Ac of steel ₁ Transformation point temperature (℃), Ac ₃ It is a predicted value of the transformation point temperature (° C.), and is a component regression equation derived by the inventors' detailed basic experiments. This predicted temperature (° C.) is particularly suitable for application when heating at a heating rate of 2 ° C./s or more and 20 ° C./s or less.
[0032]
Hereinafter, the reasons for limiting the expressions (1), (2), and (3) will be described in order.
Equation (1) is a condition for defining the addition amounts of Ti and Nb, and is based on the following knowledge.
In general, it is known that the addition of Ti and Nb precipitates TiC and NbC and has the effect of increasing the recrystallization temperature of a steel sheet. Therefore, the relationship between the added amounts of Ti and Nb and the recrystallization temperature Tre was investigated in detail. When Ti and Nb were added in a certain amount or more, the recrystallization temperature was calculated by the above equation (6). ₃ It turned out to be equivalent to
[0033]
In FIG. ₁ = 700 ° C, A ₃ The results obtained by examining the relationship between the amounts of Ti and Nb added and the recrystallization temperature Tre when the amounts of Ti and Nb are variously changed in a steel composition adjusted to = 855 ° C. Here, the recrystallization temperature Tre was determined by performing continuous annealing in a laboratory while changing the heating temperature variously, measuring the hardness, and observing the structure. The amount of Ti added is defined as the effective amount of Ti for precipitating TiC. ^* The relationship between the amounts of Ti and Nb added and the recrystallization temperature is shown using 48/93 · [% Nb] to convert the amount of Nb added to Ti.
According to the figure, 637.5 + 4930 ｛Ti ^* + (48/93) · [% Nb]} is 700 ° C, that is, A ₁ At this point, the recrystallization temperature Tre becomes around 855 ° C., ie, A ₃ It can be seen that the current rises rapidly and saturates.
[0034]
Next, FIG. 2 shows that 637.5+4930@Ti ^* + (48/93) · [% Nb]｝ ≧ A ₁ Under the conditions of A ₃ (Variation by changing C, Si, Mn, Ni, etc.) in various cases ₃ The result of examining the relationship between the temperature and the recrystallization temperature Tre is shown.
As shown in the figure, 637.5 + 4930 ｛Ti ^* + (48/93) · [% Nb]｝ ≧ A ₁ Under the conditions of the above, the recrystallization temperature Tre is A ₃ Is equivalent to
[0035]
The reason for this is not necessarily clear, but is considered as follows.
That is, Ti and Nb are added, and the recrystallization temperature rises due to the pinning force of these fine carbides. ₁ If recrystallization cannot be performed in the ferrite (α) region less than that, the temperature of the unrecrystallized α remains (ferrite + austenite (γ)) in the two-phase region. At the nucleation site, competition occurs between the generation of recrystallized α nuclei from the processed α and the generation of α → γ transformation nuclei. At this time, since the driving force of the α → γ transformation is larger than the driving force of the recrystallization, it is considered that γ nuclei are generated one after another in preference to the generation of recrystallized α nuclei and occupy the preferential nucleation site. .
Due to the rearrangement of atoms in the α → γ transformation, strain (dislocation) is consumed, and only processed α having a low dislocation density remains, and recrystallization of processed α becomes more and more difficult. The temperature rises and A ₃ , And the strain is completely eliminated only after the gamma single phase region is reached, and apparently recrystallization is completed. This is because the recrystallization temperature is A ₃ And is considered to be a mechanism of saturation.
In this case, since the α → γ transformation involves nucleation from the processed α (there are many preferential nucleation sites), the γ grains at a high temperature after recrystallization are refined. Therefore, the recrystallization temperature is set to A for the purpose of refining the high temperature γ grains during annealing. ₃ It is effective to add Ti and Nb satisfying the formula (1) in the present invention.
[0036]
Next, equation (2) is ₃ Is a condition that defines
As described above, if equation (1) is satisfied, A ₃ Is substantially at the recrystallization temperature, so A ₃ It is necessary to perform recrystallization annealing at the above temperature. Where A ₃ If the temperature exceeds 860 ° C., it is necessary to perform the recrystallization annealing at a higher temperature, and the γ grain growth is intense. As a result, fine grains having an average crystal grain size of 3.5 μm or less cannot be obtained. Therefore, A ₃ It is necessary to satisfy ≦ 860 ° C.
[0037]
Next, the expression (3) is a condition for defining the addition amount of Mn or Ni, that is, the austenite stabilizing element.
The ferrite start line in the CCT diagram shifts to the lower temperature side due to the increase of the austenite stabilizing element, and the degree of supercooling during the γ → α transformation in the cooling process after annealing increases, and α nucleates finely. Thereby, the α crystal grains are refined. Here, in order to obtain fine grains having an average crystal grain size of 3.5 μm or less, in addition to the above-described equations (1) and (2), [% Mn] + [% Ni] ≧ 1.3 ( %).
As long as [% Mn] + [% Ni] ≧ 1.3 (%) is satisfied, either Mn or Ni may be added alone or in combination. More preferably, it is in the range of [% Mn] + [% Ni] ≧ 2.0 (%).
[0038]
Next, the steel structure will be described.
In the present invention, the steel structure must have a ferrite phase volume fraction of 70 vol% or more, a retained austenite volume fraction of 5 vol% or more, and an average ferrite crystal grain size of 3.5 μm or less. is there.
That is, ferrite has a soft structure and not only contributes to improvement of ductility and workability, but also hardly forms a solid solution of carbon, and thus effectively contributes to enrichment of carbon in austenite during heat treatment before plating. Therefore, in order to obtain a hot-dip galvanized steel sheet having the strength and ductility expected in the present invention, it is necessary to have a steel structure mainly composed of fine ferrite, and in particular, a fine ferrite having an average crystal grain size of 3.5 μm or less. This is because it is important that the volume fraction of the phase is 70 vol% or more.
Here, when the average crystal grain size of ferrite exceeds 3.5 μm, the strength-elongation balance is reduced, and the stretch flangeability is deteriorated. If the volume fraction of ferrite is less than 70 vol%, ductility is significantly reduced and workability is deteriorated.
[0039]
The second phase is a retained austenite phase. The retained austenite phase is a structure that greatly contributes to the ductility of the steel sheet, and in order to obtain the effect, it is necessary that the volume fraction be 5 vol% or more of the entire structure.
[0040]
In addition, a bainite phase, a pearlite phase, and further a martensite phase may be generated in addition to the above-described ferrite phase and residual austenite phase, and if the total of these phases is 20 vol% or less in volume fraction. ,acceptable.
[0041]
Next, the manufacturing conditions will be described.
The steel adjusted to the above preferred composition is melted in a converter or the like, and is made into a slab by a continuous casting method or the like. The slab, which is a steel material, is heated to 1200 ° C. or higher after being kept in a high temperature state or cooled, and then subjected to hot rolling, and then cold-rolled. ₃ (° C) or more, (A ₃ (+30) (° C.) or lower, and then preferably cooled to 300 ° C. at a rate of 10 ° C./s or more.
[0042]
In the above process, if the heating temperature of the slab is less than 1200 ° C., TiC or the like does not sufficiently form a solid solution and becomes coarse, and the effect of raising the recrystallization temperature and the effect of inhibiting the growth of crystal grains in the subsequent recrystallization annealing step are insufficient. Therefore, the heating temperature of the slab needs to be 1200 ° C. or higher.
In the present invention, hot rolling conditions other than the slab heating temperature are not particularly limited, and may be in accordance with a conventional method. The hot finish rolling exit temperature is not particularly limited, but may be Ar ₃ Below the transformation point, α and γ are generated during rolling, and a band-like structure is easily generated in the steel sheet. Such a band-like structure remains even after cold rolling or annealing and causes anisotropy in the material properties. The hot finish rolling end temperature is Ar ₃ It is preferable that the temperature be equal to or higher than the transformation point.
[0043]
Although the winding temperature after the completion of hot rolling is not particularly limited, if it is less than 500 ° C. or more than 650 ° C., the precipitation of AlN for suppressing aging deterioration due to nitrogen is insufficient, and the material properties are poor. It will be. Further, in order to make the structure of the steel sheet uniform and to make the crystal grain size as fine and uniform as possible, it is preferable that the coil winding temperature be 500 ° C. or more and 650 ° C. or less.
[0044]
Next, preferably, the oxide scale on the surface of the hot-rolled steel sheet is removed by pickling and then subjected to cold rolling to obtain a cold-rolled steel sheet having a predetermined thickness. Here, the pickling conditions and the cold rolling conditions are not particularly limited, and may be in accordance with a conventional method.
The rolling reduction during cold rolling is desirably 40% or more from the viewpoint of increasing the number of nucleation sites during recrystallization annealing and promoting the refinement of crystal grains. Since the operation becomes difficult due to work hardening, the upper limit of the rolling reduction is preferably about 90% or less.
[0045]
Next, the obtained cold-rolled steel sheet was subjected to the temperature A shown in the above formula (6). ₃ (° C) or more, (A ₃ (+30) (° C.) or less and perform recrystallization annealing.
In the steel material of the present invention whose components have been adjusted as described above, A ₃ Is substantially equivalent to the recrystallization temperature, ₃ At a temperature lower than this, recrystallization becomes insufficient. On the other hand, (A ₃ At a temperature exceeding (+30) (° C.), the growth of γ grains during annealing is severe, which is inappropriate for miniaturization. This recrystallization annealing is preferably performed in a continuous annealing line, and the annealing time for continuous annealing is preferably about 10 seconds to 120 seconds at which recrystallization occurs. This is because recrystallization is insufficient in less than 10 seconds, and a sufficient workability may not be ensured because a processed structure that has been extended in the rolling direction and a recovered structure that has not been recrystallized remain. On the other hand, if the time is longer than 120 seconds, the γ crystal grains become coarse, and the desired strength may not be obtained.
[0046]
The rate of temperature rise during recrystallization annealing may be usually about 2 to 20 ° C./s at which continuous annealing is performed, and the cooling after recrystallization annealing is performed at an annealing temperature to suppress the growth of crystal grains. Preferably, the cooling is performed at an average cooling rate of from 10 ° C./s to 300 ° C. or higher.
[0047]
Next, after the recrystallization annealing, pickling is performed to remove surface oxides that adversely affect the plating property. That is, a surface-concentrated layer in which P, Si, Mn, Cr, and the like precipitated on the steel sheet surface during annealing are concentrated as an oxide is removed. Since such a surface-concentrated layer to be removed can be removed by light pickling, light pickling before the conventional continuous hot-dip galvanizing treatment is sufficient.
[0048]
Next, A ₁ Above, (A ₁ After heating in a temperature range of +70) ° C. or lower for 5 to 30 seconds, it is cooled preferably at a rate of 5 to 15 ° C./s to a hot-dip galvanizing starting temperature. By performing heat treatment under such conditions, substitutional alloying elements such as C, Mn, Mo, and Ni are concentrated into an austenite phase generated by transformation and stabilized, and as a result, a residual austenite phase obtained by cooling is obtained. Can be stabilized even at room temperature, and good ductility can be obtained.
[0049]
In the present invention, as described above, since the crystal grains of the steel are refined by recrystallization annealing performed prior to the heat treatment, the distance that C and other alloy elements move to the γ phase is short, and the distance to austenite is reduced. Since the alloying element such as C is easily concentrated, retained austenite can be stably obtained.
The heating temperature in this heat treatment is A ₁ If it is less than 1, an austenite phase is not generated, so that a residual austenite phase cannot be stably obtained after cooling. On the other hand, when the heating temperature is (A ₁ If the temperature exceeds (+70) ° C., crystal grain growth proceeds, and a fine structure after cooling cannot be obtained. Therefore, the heating temperature is A ₁ Above, (A ₁ +70) C or lower. On the other hand, if the heating time is less than 5 seconds, the transformation time is too short, which is disadvantageous for austenite formation. On the other hand, if the heating time is more than 30 seconds, the amount of carbon in the austenite reaches an equilibrium amount, and the effect is reduced. The amount of carbon therein decreases, which is disadvantageous in producing stable retained austenite. Therefore, the heat treatment time was limited to 5 to 30 seconds.
[0050]
Further, the cooling rate after the heat treatment is preferably set to 5 ° C./s or more in order to suppress the crystal grain growth. However, if the cooling rate is too high, martensite may be formed and ductility may be reduced. Therefore, the cooling rate is preferably set to 15 ° C./s or less.
[0051]
The above heat treatment is desirably performed in a continuous hot-dip galvanizing line.
Following the cooling after the heat treatment, hot-dip galvanizing or alloying treatment is performed to obtain a steel structure having 70 vol% or more of ferrite and 5 vol% or more of retained austenite, and the ferrite grain size is 3.5 μm. The following high-strength hot-dip galvanized cold-rolled steel sheet is obtained.
[0052]
The hot-dip galvanizing treatment or alloying treatment may be performed under ordinary conditions.
If the plating temperature is unnecessarily high, it is disadvantageous to secure the amount of retained austenite.On the other hand, if the temperature is low, the reaction between zinc and iron in the zinc bath becomes slow and the productivity decreases. The sheet temperature at this time is advantageously about 450 to 550 ° C.
[0053]
Further, in the alloying treatment after hot-dip galvanizing, if the temperature is too high, it is disadvantageous to secure the retained austenite, while if the temperature is too low, the progress of alloying is slowed down, and the productivity is lowered. Is desirably about 500 to 560 ° C.
[0054]
【Example】
A slab having the component composition shown in Table 1 was heated under the conditions shown in Table 2 and then hot-rolled according to a conventional method to obtain a hot-rolled sheet having a thickness of 4.0 mm. At this time, the hot finish rolling exit temperature is Ar ₃ The winding temperature was set to 600 ° C. or higher. The hot-rolled sheet was pickled and then cold-rolled (rolling reduction: 60%) to form a 1.6-mm-thick cold-rolled sheet, which was then annealed in a continuous annealing line under the same conditions as shown in Table 2. Then, after pickling, in a continuous hot-dip galvanizing line, a heat treatment followed by a plating treatment and an alloying treatment were performed to obtain a hot-dip galvanized steel sheet. During the cold-rolled sheet annealing, the rate of temperature rise was 5 to 10 ° C / s, and the average cooling rate up to 300 ° C was 20 to 25 ° C / s. The cooling rate from the pre-plating heating temperature to the hot-dip galvanizing temperature in the continuous hot-dip galvanizing line was 5 to 15 ° C./s. The temperature of the steel sheet in the bath in the hot-dip galvanizing treatment was 465 ° C., and the alloying temperature in the case where the alloying treatment was performed thereafter was 500 ° C. In addition, No. No. 16 was only hot-dip galvanized, and was not alloyed.
Table 3 shows the results of examining the structure and tensile properties of the hot-dip galvanized steel sheet thus obtained.
[0055]
In addition, the structure | tissue was investigated using the optical microscope or the electron microscope about the rolling direction cross section of a steel plate, and also measured the volume ratio of ferrite and the average crystal grain size. Here, the average crystal grain size of the ferrite was determined in accordance with the cutting method defined in JIS G 0552.
The tensile properties (tensile strength (TS) and elongation (EL)) were measured by a tensile test using JIS No. 5 test pieces taken from the rolling direction of the steel sheet.
The amount of retained austenite was determined by polishing a steel sheet in the thickness direction to a quarter of the thickness and measuring the X-ray diffraction intensity on this quarter plane. MoKα is used for the incident X-ray, and the {200}, {220}, and {311} faces of the retained austenite phase are compared with the X-ray diffraction intensity of the {110}, {200}, and {211} faces of the ferrite phase. X-ray diffraction intensities were determined, and the average of these values was defined as the volume fraction of retained austenite.
The stretch flangeability was evaluated by a hole expansion test described below. That is, a test piece taken according to the Japanese Iron and Steel Joint Standard JFST1001 has a diameter of 10 mmφ (D ₀ ) Is punched and then expanded with a conical punch having a vertex angle of 60 °, and the hole diameter D (mm) immediately after the crack penetrates the plate thickness is obtained.
λ = ｛(DD ₀ ) / D ₀ )｝ × 100%
Was evaluated by the hole expansion ratio λ obtained in the above.
[0056]
[Table 1]

[0057]
[Table 2]

[0058]
[Table 3]

[0059]
As shown in Table 3, in all of the invention examples, the volume fraction of ferrite is 70 vol% or more, the volume fraction of retained austenite is 5 vol% or more, and the average crystal grain size of ferrite is 3.5 μm or less. It can be seen that they are excellent in strength-ductility balance and strength-hole expansion balance in which TS × EL is 20,000 MPa ·% or more and TS × λ is 67000 MPa ·% or more.
On the other hand, no. In No. 4, TS × EL decreased because the pre-plating heating temperature greatly exceeded the appropriate upper limit temperature (732 ° C.), and crystal grain growth was intense and coarse.
No. In No. 9, since the heating temperature of the slab was low, TiC was coarsened, the effect of increasing the recrystallization temperature was suppressed, the effect of refining the crystal grains of the steel sheet was not obtained, and the crystal grain size was increased. Also, the TS × EL and TS × λ values are small.
No. In No. 10, since the annealing temperature greatly exceeded the appropriate upper limit temperature (884 ° C.) of the present invention, crystal grain growth was severe and the TS × EL value was deteriorated.
No. Sample No. 11 had an inferior TS × EL value because the annealing temperature was below the lower limit (854 ° C.) of the present invention, so that recrystallization was not completed and a processed structure remained.
No. 12 means that the heating temperature before plating is A ₁ As a result, the structure was merely tempered, an austenite phase was not formed, and the ductility was reduced.
No. 21 shows that the recrystallization occurs at a low temperature (α range) because the amounts of Ti and Nb are insufficient to satisfy the expression (1), and the recrystallization temperature and A ₃ Do not match. For this reason, during annealing after cooling, recrystallization occurred in the α range as usual, and the crystal grains became coarse, sufficient strength could not be obtained, and the elongation and hole expansion ratio were low.
No. 22 is A ₃ 860 ° C. or higher, high-temperature annealing was required, and as a result, crystal grains grew and the TS × EL value deteriorated.
No. In No. 23, since the amount of (Ni + Mn) was small, the degree of supercooling during γ-α transformation in the cooling process after annealing was small, and ferrite could not be formed into fine nuclei, so that the crystal grains became coarse.
[0060]
【The invention's effect】
Thus, according to the present invention, a high-tensile cold-rolled steel sheet having an ultrafine grain structure and excellent ductility can be stably manufactured without significant remodeling of manufacturing equipment, and is extremely industrially possible. Useful.
[Brief description of the drawings]
FIG. 1A ₁ = 700 ° C, A ₃ FIG. 4 is a diagram showing the relationship between the amounts of Ti and Nb added and the recrystallization temperature when the amounts of Ti and Nb added are variously changed in a steel composition adjusted to = 855 ° C.
FIG. 2 637.5+4930@Ti ^* + (48/93) · [% Nb]｝ ≧ A ₁ Under the conditions of A ₃ A when various values are changed ₃ FIG. 4 is a diagram showing a relationship between the temperature and a recrystallization temperature Tre.

Claims (6)

In mass%,
C: 0.03 to 0.16%,
Si: 0.2 to 2.0%,
Mn: 1.0 to 3.0% and / or Ni: 0.5 to 3.0%,
Ti: 0.2% or less and / or Nb: 0.2% or less;
Al: 0.01 to 0.1%,
P: 0.1% or less,
S: 0.02% or less and N: 0.005% or less, and C, Si, Mn, Ni, Ti and Nb are contained in the ranges satisfying the following formulas (1), (2) and (3), respectively. The remainder has a composition of Fe and unavoidable impurities. The volume fraction of ferrite is 70 vol% or more, the volume fraction of retained austenite is 5 vol% or more, and the average crystal grain size of the ferrite is 3.5 μm. A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, having the following steel structure and further having a hot-dip galvanized layer on the surface.

The steel sheet according to claim 1, wherein the steel sheet further comprises:
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that it has a composition containing one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. . The steel sheet according to claim 1 or 2, further comprising:
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that the composition contains one or more selected from Ca, REM and B in total of 0.005% or less. In mass%,
C: 0.03 to 0.16%,
Si: 0.2 to 2.0%,
Mn: 1.0 to 3.0% and / or Ni: 0.5 to 3.0%,
Ti: 0.2% or less and / or Nb: 0.2% or less;
Al: 0.01 to 0.1%,
P: 0.1% or less,
S: 0.02% or less and N: 0.005% or less, and C, Si, Mn, Ni, Ti and Nb are contained in the ranges satisfying the following formulas (1), (2) and (3), respectively. The remainder is a steel material having a composition of Fe and unavoidable impurities, heated to 1200 ° C. or more, hot-rolled, and then cold-rolled, and then subjected to a temperature A ₃ ° C. or more determined by the following equation (6). After performing recrystallization annealing at (A ₃ +30) ° C. or lower, and then pickling, heat treatment is performed for 5 to 30 seconds at a temperature of A ₁ ° C. or higher and (A ₁ +70) ° C. or lower determined by the following formula (5) A method of producing a hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, which is followed by hot-dip galvanizing treatment or further alloying treatment.

The steel material according to claim 4, further comprising:
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that it has a composition containing one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. Manufacturing method. The steel material according to claim 4 or 5, further comprising:
A hot-dip galvanized steel sheet having an ultrafine grain structure and excellent ductility, characterized in that the composition contains one or more selected from Ca, REM and B in a total content of 0.005% or less. Production method.

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