JP2004211138A - Galvanized cold-rolled steel sheet having ultrafine granular structure and excellent fatigue characteristic, and its producing method - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high tension cold-rolled steel sheet having ultrafine granular structure and especially, excellent in the balance of strength-elongation and fatigue characteristic in the mechanical characteristics. <P>SOLUTION: This steel has a composition of especially, C, Si, Mn, Ni, Ti and Nb in the respective ranges satisfying the following formulas (1), (2), (3) and the balance Fe with inevitable impurities in the steel components, and has the structure of ≥65 vol% ferrite having ≤3.5 μm average crystal grain diameter and 10 vol% martensite phase as the second phase. 637.5+4930äTi<SP>*</SP>+(48/93)×[%Nb]}≤A<SB>1</SB>: (1), A<SB>3</SB>≤860: (2), [%Mn]+[%Ni]≥1.3: (3). Wherein, Ti<SP>*</SP>=[%Ti]-(48/32)×[%S]-(48/14)×[%N], A<SB>1</SB>: predicted value (°C) of transformation point A<SB>1</SB>obtained with a calculating equation, A<SB>3</SB>: predicted value (°C) of transformation point A<SB>3</SB>obtained with a calculating equation. <P>COPYRIGHT: (C)2004,JPO&NCIPI

Description

また、［％Ｍ］ はＭ元素の含有量（質量％）
【００１３】
２．上記１において、鋼板が、さらに質量％で、
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疲労特性に優れる溶融亜鉛めっき冷延鋼板。
【００１４】
３．上記１または２において、鋼板が、さらに質量％で、
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
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疲労特性に優れる溶融亜鉛めっき冷延鋼板。
【００１５】
４．質量％で、
Ｃ：０．０３〜０．１６％、
Ｓｉ：２．０ ％以下、
Ｍｎ：３．０ ％以下および／またはＮｉ：３．０ ％以下、
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、熱間圧延を行い、巻取り後、冷間圧延を施したのち、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、その後 ６００℃まで５℃／ｓ以上の速度で冷却し、ついで酸洗後、下記（５）式で求められる（Ａ_１ ＋７０）（℃）以上、Ａ_３ （℃） 以下の温度範囲で熱処理を行い、その後溶融亜鉛めっき処理温度まで５℃／ｓ以上の速度で冷却し、溶融亜鉛めっきを施したのち、３００ ℃まで５℃／ｓ以上の速度で冷却するか、上記の溶融亜鉛めっき後、さらに合金化処理を施したのち、３００ ℃まで５℃／ｓ以上の速度で冷却することを特徴とする、超微細粒組織を有し疲労特性に優れる溶融亜鉛めっき冷延鋼板の製造方法。

また、［％Ｍ］ はＭ元素の含有量（質量％）
【００１６】
５．上記４において、巻取り温度が ７５０℃以下、 ６５０℃以上であることを特徴とする、超微細粒組織を有し疲労特性に優れる溶融亜鉛めっき冷延鋼板の製造方法。
【００１７】
６．上記４または５において、鋼素材が、さらに質量％で、
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疲労特性に優れる溶融亜鉛めっき冷延鋼板の製造方法。
【００１８】
７．上記４，５または６において、鋼素材が、さらに質量％で、
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
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疲労特性に優れる溶融亜鉛めっき冷延鋼板の製造方法。
【００１９】
【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、マルテンサイトを生成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延性や溶接性が劣化するため、Ｃは０．０３〜０．１６％の範囲に限定した。
【００２０】
Ｓｉ：２．０ ％以下
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を向上させるのに有効に寄与するが、過剰な添加は、延性や表面性状、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。なお、好ましくは０．０１〜０．６ ％の範囲である。
【００２１】
Ｍｎ：３．０ ％以下および／またはＮｉ：３．０ ％以下
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、多量の添加は鋼を硬質化し、却って強度−延性バランスを劣化させるので、いずれも ３．０％以下で含有させるものとした。
なお、Ｍｎは、有害な固溶ＳをＭｎＳとして無害化する作用も併せて有するので、０．１ ％以上含有させることが好ましい。また、Ｎｉは０．０１％以上含有させることが好ましい。
【００２２】
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。そのためには、それぞれ０．０１％以上含有させることが好ましい。そして、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。
【００２３】
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。
【００２４】
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効な元素であるが、一方で多量の含有は加工性や靱性の低下を招くので、Ｐは ０．１％以下で含有させるものとした。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。
【００２５】
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伸びフランジ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。
【００２６】
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。
【００２７】
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。
【００２８】
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０．０００５％以上含有させることが好ましい。
【００２９】
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。

また、［％Ｍ］ はＭ元素の含有量（質量％）
【００３０】
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から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。
【００３１】
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。
【００３２】
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。
【００３３】
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。
【００３４】
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 に調整することは極めて有効であるので、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。
【００３５】
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。なお好ましくはＡ_３ ≦ ８３０℃である。
【００３６】
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。
【００３７】
次に、鋼組織について説明する。
本発明では、鋼組織は、フェライト相の組織分率を体積率で６５％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。
というのは、本発明で所期した強度、延性、靱性および強度−伸びバランスに優れた冷延鋼板とするには、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組織分率を６５ ｖｏｌ％以上とすることが重要だからである。より好ましくは７５ ｖｏｌ％以上である。
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−伸びバランスが劣化すると共に、靱性が低下し、また軟質なフェライトの組織分率が６５ ｖｏｌ％に満たないと延性が著しく低下し、加工性に乏しくなる。
【００３８】
また、フェライト以外の第２相組織は、マルテンサイトとする必要がある。
というのは、第２相を硬質なマルテンサイトとすることにより、強度−伸びバランスを向上せしめ、かつ疲労特性が向上するからである。
なお、疲労特性が向上するメカニズムは必ずしも明らかではないが、硬質第２相の存在により、進展する疲労亀裂が偏向し、結果として亀裂伝播の抑制効果が作用しているものと考えられる。
【００３９】
なお、フェライトおよびマルテンサイト以外の比較的軟質なベイナイトや延性に劣るパーライト等の相が多量に存在すると、強度−伸びバランスが劣化すると共に、先に述べたように硬質第２相の存在による疲労亀裂伝播の抑制効果を得難く、疲労特性が劣化するおそれが生じるが、これらの相が体積分率で３％未満であれば、許容できる。
【００４０】
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。この鋼素材を、高温状態のまま、あるいは一旦冷却したのち、１２００℃以上に加熱してから、熱間圧延を行い、ついで好適には ７５０℃以下、 ６５０℃以上の温度でコイルに巻取り、好ましくは酸洗後、冷間圧延を施した後に、温度Ａ_３（℃）以上、（Ａ_３ ＋３０）（℃）以下で再結晶焼鈍を施し、その後少なくとも ６００℃まで５℃／ｓ以上の速度で冷却する。
【００４１】
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。
【００４２】
なお、本発明において、熱間圧延時の熱間圧延終了温度が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、通常よく行われているように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】
また、熱延終了後の巻取り過程においては、 ７５０〜６５０ ℃の温度域まで冷却して巻取り、この間にγ→α変態を生じさせると共に、熱延板にＴｉＣやＮｂＣを微細に析出させることが好ましい。熱延板の段階で、このような微細析出物を析出させることは、製品板において、フェライト相を硬質化し、フェライト相自体の疲労特性を改善し、ひいては製品板の疲労特性をさらに改善できるので、好ましい。ここに、巻取り温度が ６５０℃未満では、Ｔｉ，Ｎｂ等の炭化物を主体とした微細析出物が得難く、一方 ７５０℃を超えると、パーライト変態が進んでＴｉやＮｂ等の炭化物を主体とした微細析出物生成に必要な炭素が消費されるため、やはり微細な析出物が得難くなる。
【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となり易く、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】
引き続き、焼鈍温度から少なくとも ６００℃まで、冷却速度：５℃／ｓ以上の条件で冷却する。なお、ここで冷却速度は、焼鈍温度から ６００℃までの平均冷却速度である。ここに、上記冷却速度が５℃／ｓ未満では、冷却中におけるγ→α変態時の過冷度が小さく、結晶粒径が粗大化する。よって、焼鈍温度から ６００℃までの冷却速度は５℃／ｓ以上とする必要がある。
また、上記の制御冷却処理の終点温度を ６００℃としたのは、結晶粒の微細化にはγ→α変態が開始する ６００℃までが強く影響するからである。
上記までの製造工程により、フェライト平均粒径が ３．５μｍ 以下の超微細組織を得ることができる。
【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、再結晶焼鈍中に鋼板表面に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前の軽酸洗で十分である。
【００４８】
つぎに、前掲した（５）式で求められる（Ａ_１ ＋７０）（℃）以上、（６）式で求められるＡ_３ （℃） 以下の温度範囲で１０秒以上の熱処理を行い、溶融亜鉛めっき処理温度まで５ ℃／ｓ以上の速度で冷却する。
この熱処理により，（α＋γ）２相域に加熱し、γ相中にＣおよび合金元素を濃化させて焼入れ性を向上させ、比較的緩やかな冷却条件でも硬質なマルテンサイトを得易くして強度−延性バランスを確保すると共に，疲労特性の改善を図るのである。
【００４９】
なお、本発明では、先行して行う再結晶焼鈍で結晶粒を微細化している関係上、ＣやＭｎなどの他の合金元素がγ相へと移動する距離が短く、γ相へのＣなどの合金元素が濃化し易いため、焼入れ性が向上してマルテンサイト相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_３ （℃） 超えでは、完全オーステナイト相となるため２相分離せずγ中への成分濃化が進行せず、冷却後に十分なマルテンサイト相が得られないため，良好な疲労特性が得難く、また結晶粒の成長も余儀なくされ、強度−延性バランスも劣化する。一方、加熱温度が（Ａ_１ ＋７０）（℃）未満では、加熱中のγ分率が低いため，結果として得られるマルテンサイト相が十分ではなく，やはり良好な疲労特性は得られない。よって，加熱温度は（Ａ_１ ＋７０）（℃）を以上、Ａ_３ （℃） 以下の範囲に限定した。
【００５０】
また、上記加熱時間が１０秒未満では、変態できる時間が短くオーステナイト相へのＣやＭｎなどの他の合金元素の濃化すなわちマルテンサイトの生成に不利となるため、１０秒以上とすることが好ましい。なお，実操業上の効率を考慮すると、加熱時間の上限は １２０秒程度とすることが好ましい。
【００５１】
さらに、熱処理後の冷却速度は，γ相のマルテンサイト変態の促進および結晶粒成長抑止の観点から、溶融亜鉛めっき処理温度まで５ ℃／ｓ以上とする必要がある。
なお、上記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。
【００５２】
上記しためっき前熱処理に引き続き、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、平均結晶粒径が ３．５μｍ 以下のフェライトが主相で、第２相として硬質なマルテンサイトを有する組織の高張力溶融亜鉛めっき冷延鋼板を得ることができる。
【００５３】
本発明における溶融亜鉛めっき処理は、通常の溶融亜鉛めっきラインで行われる条件と同様に ４５０〜５５０ ℃の温度範囲で溶融亜鉛めっきを施し、鋼板表面に溶融亜鉛めっき層を形成する。
また、本発明では、溶融亜鉛めっき処理後、溶融亜鉛めっき層を合金化する合金化処理を施すこともできる。この合金化処理における処理温度は ４７０〜５７０ ℃程度とするのが好ましい。というのは、加熱温度が ４７０℃未満では合金化の進行が遅く生産性の低下を招き、一方 ５７０℃を超えるとめっき層の合金化が進行しすぎて合金化溶融亜鉛めっき層が脆化するからである。
【００５４】
さらに、めっき後に必要なマルテンサイト量を確保し、粒成長を抑制するために、めっき後 ３００℃までの冷却速度、あるいはさらに合金化処理を行う場合は、合金化処理後 ３００℃までの冷却速度は５℃／ｓ以上とする必要がある。
かくして、上記の製造方法とすることにより、超微細粒組織を有し、強度−延性バランスに優れ、さらには疲労特性に優れる溶融亜鉛めっき冷延鋼板を得ることができる。
【００５５】
【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。なお，仕上げ圧延温度は、常法に従い全てＡｒ_３点以上とした。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に示す条件下で再結晶焼鈍を行い、ついで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行って、製品板とした。この時、冷延板焼鈍時の昇温速度は５〜１０℃／ｓとした。また、溶融亜鉛めっき処理における浴温は ４６５℃、その後合金化処理を施した場合の合金化温度は ５２０℃とした。なお、Ｎｏ．１９ は溶融亜鉛めっきのみとし合金化処理は施さなかった。
かくして得られた製品板の組織、引張特性および疲労特性について調査した結果を表３に示す。
【００５６】
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観察し、フェライトの平均結晶粒径を求めると共に、各組識の面積率を求めてこれを体積率とした。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収したＪＩＳ５号試験片を用いた引張試験により測定した。
さらに、疲労特性は、図３に示す寸法形状になるＪＩＳ Ｚ ２２７５（金属平板の平板曲げ疲れ試験方法）の試験片を用いて、両振りの繰り返し曲げ試験により求めた。この時、１０００万サイクルに到達した時点を疲労限ＦＬとした。
【００５７】
【表１】

【００５８】
【表２】

【００５９】
【表３】

【００６０】
表３に示したとおり、発明例はいずれも、主相すなわち６５％以上の分率を占めるフェライトの平均粒径が ３．５μｍ 以下と微細であり、特にＮｉ，Ｍｎ量を増量し、Ａ_３ を低下させたＧ鋼を用いた Ｎｏ．１６は、平均結晶粒径が ０．９μｍ と超微細粒となっている。また，発明例はいずれも、ＴＳ×ＥＬが １７５００ ＭＰａ・％以上と強度−延性バランスに優れており、さらに疲労特性については，疲労限ＦＬとＴＳの比（ＦＬ／ＴＳ）が０．４５以上と高いレベルが得られている。これは第２相としての硬質マルテンサイトの存在、さらには結晶粒微細化の効果により、疲労亀裂の発生と亀裂伝播が抑制されたためと考えられる。また、巻取り温度を前述の好適範囲内とすることにより、フェライト相がＴｉＣやＮｂＣ等の微細析出物で硬質化できたためと考えられるが、さらに疲労特性を向上させることができた。
【００６１】
これに対し、Ｎｏ．５は、めっき前加熱温度が低く、加熱中のオーステナイト分率が低かったため、結果としてマルテンサイト分率が低く、ＴＳ×ＥＬ、ＦＬ／ＴＳ 値も小さくなっている。
Ｎｏ．６は、逆にめっき前加熱温度が高すぎたため、加熱中の２相分離によるγへの成分濃化が進行せず、やはりマルテンサイト分率が低く、ＦＬ／ＴＳ 値の低下を招いた。
Ｎｏ．７は、めっき前加熱後めっき処理温度までの冷却速度および合金化処理後の冷却速度が小さかったため、第２相がマルテンサイトよりも軟質なべイナイト主体となり、ＴＳ×ＥＬ、ＦＬ／ＴＳ 値が劣った。
Ｎｏ．１１ 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径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、ＦＬ／ＴＳ 値も低下している。
Ｎｏ．１２ は、再結晶焼鈍温度が本発明の適正上限温度（８４６ ℃）を大きく超えたため、結晶粒成長が激しく、ＴＳ×ＥＬ、ＦＬ／ＴＳ が劣化した。
Ｎｏ．１３ は、焼鈍温度が本発明の下限（ ８１６℃）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値が極めて劣っている。
Ｎｏ．１４ は、再結晶焼鈍後の冷却速度が小さかったために、結晶粒が粗大化して強度が低下し、ＴＳ×ＥＬ値の劣化を招いた。
Ｎｏ．２３ は、Ｔ_Ｘ がＡ_１ 未満であることから、再結晶焼鈍によるγ粒微細化効果が得られず、粗大粒となったため強度の低下をきたし、その結果ＴＳ×ＥＬ、ＦＬ／ＴＳ値の低下を招いた。
Ｎｏ．２４ は、Ａ_３ が ８６０℃を超えていることから、高温焼鈍が必要となり、その結果結晶粒が成長して、ＴＳ×ＥＬ値の低下を招いた。
Ｎｏ．２５ 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、α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。
【００６２】
【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、機械的特性なかでも強度−伸びバランスに優れ、さらには疲労特性にも優れた高張力溶融亜鉛めっき冷延鋼板を、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。
【図３】平板曲げ疲れ試験用の試験片の寸法形状を示した図である。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention provides a hot-dip galvanized cold-rolled steel sheet suitable for use as a steel for automobiles and home appliances, and even mechanical structures, particularly a high-strength hot-dip galvanized steel sheet having an ultrafine grain structure and excellent strength, ductility and fatigue properties. The present invention relates to a cold-rolled steel sheet and a method for producing the same.
[0002]
[Prior art]
BACKGROUND ART Steel sheets used for automobiles, home appliances, and steel sheets for machine structures generally use various surface-treated steel sheets because they are required to have good workability and corrosion resistance. Among them, the hot-dip galvanized steel sheet has not only high corrosion resistance, but also the advantage that it can be manufactured at very low cost by a continuous hot-dip galvanizing line (CGL) that can perform recrystallization annealing and galvanizing on the same line. It has. Further, a so-called alloyed hot-dip galvanized steel sheet which has been heated and alloyed immediately after the hot-dip galvanizing, is particularly excellent in corrosion resistance.
[0003]
On the other hand, steel materials used as steel plates for automobiles, home appliances and mechanical structures are required to have not only mechanical properties such as strength and workability but also excellent fatigue properties (also called durability). Since it is effective to refine the structure as a means for comprehensively improving such properties, many production methods for obtaining a fine structure have been proposed so far.
[0004]
As a means for refining the structure, a large rolling reduction method has been conventionally known. The key point of the mechanism for refining the structure in the large rolling reduction method is to apply a large reduction to austenite grains to promote γ-α strain-induced transformation (for example, see Patent Documents 1 and 2).
Also, a case where a controlled rolling method or a controlled cooling method is applied is known (for example, see Patent Document 3).
[0005]
In addition, for the base steel, a steel structure at least partially composed of ferrite is set, and while being subjected to plastic working, the temperature is raised to a temperature range above the transformation point (Ac ₁ point), or _After maintaining at _one or more temperature ranges for a certain period of time, a part or the whole of the structure is once transformed back into austenite, ultra-fine austenite grains appear, and then cooled to an average grain size of 5 μm or less. A technique for forming a structure mainly composed of anisotropic ferrite grains has been proposed (for example, see Patent Document 4).
[0006]
On the other hand, as means for improving the workability and fatigue properties required for high-strength steel sheets, control of the existence and precipitation of a hard second phase (mainly martensite) that plays a role in inhibiting the propagation of fatigue cracks is considered. It is also known to carry out such a technique, and a technique for obtaining a steel sheet having excellent fatigue properties and workability by applying both precipitation strengthening and structure strengthening to a hot-rolled steel sheet has been proposed (for example, see Patent Document 5).
According to this technique, the hard second phase suppresses crack propagation (which is considered to be a deflection effect of fatigue crack growth due to the presence of the hard second phase), thereby improving fatigue characteristics, and at the same time, the precipitate forms a soft ferrite phase. It is described that the hardness difference between the second phase and the ferrite is reduced by the strengthening, so that the deformed portions are dispersed, so that the hole expanding property, that is, the stretch flangeability is improved.
[0007]
All of the above techniques are techniques for making crystal grains fine in a hot rolling process, that is, techniques for making hot rolled sheets finer.
In order to reduce the body weight of automobiles in consideration of environmental issues, it is effective to apply high-strength steel and reduce the thickness. Since an added alloy element increases, a larger rolling load is generally required, and it becomes difficult to produce a thin hot-rolled steel sheet. For such manufacturing reasons, cold-rolled steel sheets and hot-dip galvanized steel sheets having corrosion resistance are in high demand for high-strength thin materials.
However, with respect to cold-rolled steel sheets, there is almost no technique for making crystal grains fine in a normal cold rolling-annealing process.
[0008]
[Patent Document 1] JP-A-53-123823 (Claims)
[Patent Document 2] Japanese Patent Publication No. 5-65564 (Claims)
[Patent Document 3] JP-A-63-128117 (Claims)
[Patent Document 4] JP-A-2-301540 (Claims)
[Patent Document 5] JP-A-5-179396 (Claims)
[0009]
[Problems to be solved by the invention]
The present invention has been developed in view of the above-mentioned current situation, and enables a hot-dip galvanized steel sheet having excellent corrosion resistance to be used as a steel sheet for automobiles, home appliances, and mechanical structures, to make fine graining possible, and to reduce fatigue. It is an object of the present invention to propose a hot-dip galvanized cold-rolled steel sheet having improved properties together with its advantageous production method.
[0010]
Here, the target values of the strength, workability and fatigue properties of the hot-dip galvanized cold-rolled steel sheet in the present invention are as follows.
・ Tensile strength (TS) ≧ 590 MPa
・ Strength-elongation balance (TS × El) ≧ 17500 MPa ·%
-Ratio FL / TS ≧ 0.45 between fatigue limit FL and tensile strength TS
[0011]
[Means for Solving the Problems]
By the way, the present inventors have conducted intensive studies on the hot-dip galvanized cold-rolled steel sheet in order to achieve ultra-fine graining and achieve the above-mentioned object. As a result, the alloying elements were appropriately adjusted, and the recrystallization temperature of the steel sheet and A ₁ and after having controlled the a ₃ transformation temperature, by optimizing the recrystallization annealing temperature and subsequent cooling rate after cold rolling, the average crystal grain size is 3.5 [mu] m or less of ultrafine grained structure obtained, further By optimizing the heat history required for hot-dip galvanizing, the second phase can be optimized, and thus it has been found that not only the strength-ductility balance but also the fatigue properties are remarkably improved. .
Further, it has been found that by further optimizing the winding temperature, the fatigue characteristics are further improved.
The present invention is based on the above findings.
[0012]
That is, the gist configuration of the present invention is as follows.
1. In mass%,
C: 0.03 to 0.16%,
Si: 2.0% or less,
Mn: 3.0% or less and / or Ni: 3.0% or less,
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, and has a structure in which ferrite having an average crystal grain size of 3.5 μm or less is 65 vol% or more, and martensite is 10 vol% or more as a second phase, Furthermore, a hot-dip galvanized cold-rolled steel sheet having an ultra-fine grain structure and excellent fatigue properties, characterized by having a hot-dip galvanized layer on the surface.

[% M] is the content of element M (% by mass).
[0013]
2. In the above item 1, the steel sheet may further contain
Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. Cold rolled steel sheet.
[0014]
3. In the above 1 or 2, the steel sheet further contains, by mass%,
Hot-dip galvanizing cold-rolling having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains a total of 0.005% or less of one or more selected from Ca, REM and B. steel sheet.
[0015]
4. In mass%,
C: 0.03 to 0.16%,
Si: 2.0% or less,
Mn: 3.0% or less and / or Ni: 3.0% or less,
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 obtained by heating a steel material having a composition of Fe and unavoidable impurities to 1200 ° C. or higher, performing hot rolling, winding, and then performing cold rolling, and then obtaining the following equation (6). The recrystallization annealing is performed at a temperature of not less than A ₃ (° C.) and not more than (A ₃ +30) (° C.), and then cooled to 600 ° C. at a rate of 5 ° C./s or more. Heat treatment is performed at a temperature in the range of (A ₁ +70) (° C.) or more and A ₃ (° C.) or less, and then cooled at a rate of 5 ° C./s or more to the hot-dip galvanizing treatment temperature to perform hot-dip galvanizing. After that, cool to 300 ° C at a rate of 5 ° C / s or more. Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that after the above-mentioned hot-dip galvanizing, the alloying treatment is further performed and then cooled to 300 ° C. at a rate of 5 ° C./s or more. Manufacturing method of cold rolled steel sheet.

[% M] is the content of element M (% by mass).
[0016]
5. 4. The method for producing a hot-dip galvanized cold-rolled steel sheet according to 4 above, wherein the coiling temperature is 750 ° C. or lower and 650 ° C. or higher, and the steel sheet has an ultrafine grain structure and excellent fatigue properties.
[0017]
6. In the above 4 or 5, the steel material further contains, by mass%,
Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. Manufacturing method of cold rolled steel sheet.
[0018]
7. In the above 4, 5 or 6, the steel material further contains, by mass%,
Hot-dip galvanizing cold-rolling having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains a total of 0.005% or less of one or more selected from Ca, REM and B. Steel plate manufacturing method.
[0019]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described specifically.
First, the reason why the composition of steel 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 reinforcing component, but also a useful element in forming martensite. However, when the content is less than 0.03%, the effect of the addition is poor. On the other hand, when the content exceeds 0.16%, ductility and weldability deteriorate, so C is 0.03 to 0.16%. Limited to the range.
[0020]
Si: 2.0% or less Si, as a solid solution strengthening component, effectively contributes to improving strength while improving strength-elongation balance, but excessive addition deteriorates ductility, surface properties, and weldability. Therefore, the content of Si was set to 2.0% or less. Incidentally, the content is preferably in the range of 0.01 to 0.6%.
[0021]
Mn: 3.0% or less and / or Ni: 3.0% or less Both Mn and Ni are austenite stabilizing elements and contribute to the refinement of crystal grains through the action of lowering the A ₁ and A ₃ transformation points. In addition, it has the effect of enhancing the strength-ductility balance through the effect of promoting the formation of the second phase. However, the addition of a large amount hardens the steel and rather degrades the strength-ductility balance, so that the content of each was set to 3.0% or less.
Since Mn also has a function of detoxifying harmful solute S as MnS, it is preferable to contain Mn in an amount of 0.1% or more. Further, it is preferable that Ni is contained at 0.01% or more.
[0022]
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 the recrystallization temperature of the steel sheet is increased. For that purpose, it is preferable to contain each in 0.01% or more. Each of these may be added alone or in combination, but adding any of them in excess of 0.2% not only saturates the effect, but also increases the amount of precipitates and increases the amount of ferrite. Therefore, the content of each of them is set to 0.2% or less.
[0023]
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.
[0024]
P: 0.1% or less P is an element effective in achieving high strength at low cost without causing a large decrease in ductility, but a large amount of P causes a decrease in workability and toughness. , P are contained at 0.1% or less. In the case where demands on workability and toughness are severe, it is preferable to reduce P. In this case, it is preferable that P is set to 0.02% or less.
[0025]
S: not more than 0.02% S not only causes hot cracking during hot rolling, but also exists as inclusions such as MnS in the steel sheet and causes deterioration of ductility and stretch flangeability. Is desirable, but up to 0.02% can be tolerated.
[0026]
N: 0.005% or less Nitrogen not only causes aging deterioration but also causes yield elongation, so that it is controlled to 0.005% or less.
[0027]
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 types of Mo and Cr selected from Mo: 1.0% or less and Cr: 1.0% or less can be contained as a strengthening component, if necessary. On the contrary, since the strength-ductility balance is deteriorated, it is desirable to contain each in an amount of 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.
[0028]
One or more selected from Ca, REM and B in total 0.005% or less Ca, REM and B all have the effect of improving workability by controlling sulfide morphology and increasing 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 exhibit the above-described effects, it is preferable that one or more selected from Ca, REM, and B are contained in an amount of 0.0005% or more.
[0029]
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.

[% M] is the content of element M (% by mass).
[0030]
The above A ₁ and A ₃ are predicted values of the Ac ₁ transformation point temperature (° C.) and the Ac ₃ transformation point temperature (° C.) of the steel, respectively, and are components derived from the inventors' detailed basic experiments. It is a regression equation. 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.
[0031]
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, Ti, was investigated in detail the relationship of the added amount of Nb and recrystallization temperature Tre, Ti, is added over a certain amount of Nb, the recrystallization temperature becomes equivalent to A ₃ which is calculated by the equation (6) It has been found.
[0032]
FIG. 1 shows the relationship between the recrystallization temperature Tre and the amounts of Ti and Nb added when the amounts of Ti and Nb were variously changed in steel compositions adjusted to A ₁ = 700 ° C. and A ₃ = 855 ° C. The results are shown. 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 was determined by using Ti ^* as an effective Ti amount for precipitating TiC, and the amount of Nb added was calculated by using 48/93 · [% Nb] to convert into Ti. It shows the relationship with the crystallization temperature.
According to the ^{figure, 637.5 + 4930 {Ti * +} (48/93) · [% Nb]} When is 700 ° C. That is _{A 1} or more, the recrystallization temperature Tre is soared near 855 ° C. vicinity i.e. _{A 3} saturated You can see that
[0033]
Next, variation in FIG. 2, in ^{637.5 + 4930 {Ti * + (} 48/93) · [% Nb]} conditions _{≧ A 1, A 3 (C} , Si, Mn, by changing the Ni or the like ) shows the results of examining the relationship between the recrystallization temperature Tre and a ₃ in the case where a was variously changed.
As shown in the drawing, on the ^{637.5 + 4930 {Ti * + (} 48/93) · [% Nb]} conditions ≧ _{A 1,} the recrystallization temperature Tre has become equivalent to _{A 3.}
[0034]
The reason for this is not necessarily clear, but is considered as follows.
That, Ti, Nb are added, recrystallization temperature is increased by the pinning force of those fine carbides, when it becomes impossible to re-crystallization by A ₁ below the ferrite (alpha) region, remains machining alpha of non-recrystallized (Ferrite + austenite (γ)) When the temperature reaches the two-phase region, competition of recrystallization α nucleation from processing α and α → γ transformation nucleation at preferential nucleation sites such as high dislocation density and uneven deformation Occurs. 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 the γ nuclei are generated one after another in preference to the generation of the 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. Temperature rises greater than A _3, is eliminated completely for the first time distortion becomes γ single phase region is apparently recrystallization is completed. This recrystallization temperature matches the A _3, is believed mechanism is saturated.
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, since adjusting the recrystallization temperature to the A ₃ for high temperature γ grains finer during annealing is very effective, Ti in the present invention satisfying formula (1), since the addition of Nb is there.
[0035]
Next, (2) is a condition for defining the _{A 3.}
As described above, in the case of satisfying the expression (1) is, A ₃ is to become substantially the recrystallization temperature, it is necessary to perform recrystallization annealing at A ₃ or higher. Here, when A ₃ exceeds 860 ° C., it is necessary to perform the recrystallization annealing at a higher temperature, and γ grain growth is intense. As a result, fine grains having an average crystal grain size of 3.5 μm or less are obtained. Did not. Therefore, it is necessary to satisfy A ₃ ≦ 860 ° C. It is more preferable that A ₃ ≦ 830 ° C.
[0036]
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 (%).
[0037]
Next, the steel structure will be described.
In the present invention, in the steel structure, the structure fraction of the ferrite phase is 65% or more by volume, and the average crystal grain size of the ferrite is 3.5 μm or less.
That is, in order to obtain a cold-rolled steel sheet excellent in strength, ductility, toughness and strength-elongation balance expected in the present invention, it is necessary to have a steel structure mainly composed of fine ferrite, and particularly, an average crystal grain size. This is because it is important that the microstructure of the fine ferrite phase having a particle size of 3.5 μm or less be 65 vol% or more. More preferably, it is at least 75 vol%.
Here, if the average crystal grain size of ferrite exceeds 3.5 μm, the strength-elongation balance is deteriorated, and the toughness is reduced. If the structure fraction of soft ferrite is less than 65 vol%, the ductility is significantly reduced. And poor workability.
[0038]
Further, the second phase structure other than ferrite needs to be martensite.
This is because by making the second phase hard martensite, the strength-elongation balance is improved and the fatigue properties are improved.
The mechanism by which the fatigue properties are improved is not necessarily clear, but it is considered that the propagation of the fatigue crack is deflected by the presence of the hard second phase, and as a result, the effect of suppressing the crack propagation is acting.
[0039]
In addition, if there are a lot of relatively soft phases other than ferrite and martensite, such as bainite and inferior ductility, such as pearlite, the strength-elongation balance is deteriorated, and as described above, the fatigue due to the presence of the hard second phase is reduced. It is difficult to obtain the effect of suppressing crack propagation, and there is a possibility that fatigue properties may be deteriorated. However, if these phases are less than 3% by volume fraction, it is acceptable.
[0040]
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. This steel material is kept at a high temperature or once cooled and then heated to 1200 ° C. or more, then hot-rolled, and then preferably wound around a coil at a temperature of 750 ° C. or less and 650 ° C. or more, Preferably, after pickling and cold rolling, recrystallization annealing is performed at a temperature of A ₃ (° C.) or more and (A ₃ +30) (° C.) or less, and then a rate of 5 ° C./s or more up to at least 600 ° C. Cool with.
[0041]
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.
[0042]
In the present invention, if the hot rolling end temperature during hot rolling is lower than the Ar ₃ transformation point, α and γ occur during rolling, and a band-like structure is easily generated on the steel sheet. Since it may remain after cold rolling or after annealing and cause anisotropy in the material properties, the finish rolling end temperature is usually set to the Ar ₃ transformation point or higher as is usually performed. preferable.
[0043]
In the winding process after the completion of hot rolling, cooling is performed to a temperature range of 750 to 650 ° C. and winding is performed, during which γ → α transformation is caused and TiC or NbC is finely precipitated on the hot rolled sheet. Is preferred. Precipitating such fine precipitates at the stage of hot-rolled sheet can harden the ferrite phase in the product sheet, improve the fatigue properties of the ferrite phase itself, and further improve the fatigue properties of the product sheet. ,preferable. Here, if the winding temperature is lower than 650 ° C., it is difficult to obtain fine precipitates mainly composed of carbides such as Ti and Nb, while if it exceeds 750 ° C., the pearlite transformation proceeds, and carbides such as Ti and Nb are mainly formed. Since carbon required for producing the fine precipitates is consumed, it is difficult to obtain fine precipitates.
[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 is heated to a temperature not less than A ₃ (° C.) and not more than (A ₃ +30) (° C.) shown in the above formula (6) to perform recrystallization annealing.
The steel material of the present invention components adjusted as described above, the A ₃ is in the recrystallization temperature equivalent, recrystallization is insufficient at temperatures below A _3. On the other hand, at a temperature exceeding (A ₃ +30) (° C.), γ grains grow rapidly during annealing, 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 likely to be insufficient in a time shorter 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]
Subsequently, cooling is performed from the annealing temperature to at least 600 ° C. at a cooling rate of 5 ° C./s or more. Here, the cooling rate is an average cooling rate from the annealing temperature to 600 ° C. If the cooling rate is less than 5 ° C./s, the degree of supercooling during the γ → α transformation during cooling is small, and the crystal grain size becomes coarse. Therefore, the cooling rate from the annealing temperature to 600 ° C. needs to be 5 ° C./s or more.
In addition, the reason why the end point temperature of the above-mentioned controlled cooling treatment is set to 600 ° C. is that up to 600 ° C. at which the γ → α transformation starts is strongly affected in refining the crystal grains.
By the above manufacturing steps, an ultrafine structure having an average ferrite grain size of 3.5 μm or less can be obtained.
[0047]
Next, after the recrystallization annealing, pickling is performed to remove surface oxides that adversely affect the plating property. That is, during the recrystallization annealing, a surface-concentrated layer in which P, Si, Mn, Cr, and the like are concentrated as oxides on the steel sheet surface is removed. It should be noted that such a surface-concentrated layer to be removed can be removed by light pickling, so that conventional light pickling before continuous galvanizing is sufficient.
[0048]
Next, a heat treatment is performed for 10 seconds or more in a temperature range not less than (A ₁ +70) (° C.) obtained by the above-described equation (5) and not more than A ₃ (° C.) obtained by the above-mentioned equation (6), and the hot-dip galvanizing is performed. Cool at 5 ° C / s or more to the processing temperature.
By this heat treatment, the steel is heated to the (α + γ) two phase region, C and alloy elements are concentrated in the γ phase to improve the hardenability, and hard martensite can be easily obtained even under relatively moderate cooling conditions. -To improve the fatigue properties while maintaining the ductility balance.
[0049]
In the present invention, the distance that other alloying elements such as C and Mn move to the γ phase is short due to the refining of the crystal grains by the preceding recrystallization annealing. Since the alloying element is easily concentrated, the quenchability is improved and the martensite phase can be obtained stably. If the heating temperature in this heat treatment exceeds A ₃ (° C.), a complete austenite phase is not formed, and two phases are not separated, and the concentration of the components in γ does not proceed. After cooling, a sufficient martensite phase cannot be obtained. It is difficult to obtain good fatigue characteristics, the crystal grains must be grown, and the strength-ductility balance is deteriorated. On the other hand, when the heating temperature is less than (A ₁ +70) (° C.), the resulting martensite phase is not sufficient because the γ fraction during heating is low, and good fatigue properties cannot be obtained. Therefore, the heating temperature was limited to the range from (A ₁ +70) (° C.) to A ₃ (° C.).
[0050]
If the heating time is less than 10 seconds, the transformation time is short, which is disadvantageous for enrichment of other alloying elements such as C and Mn into the austenite phase, that is, the formation of martensite. preferable. In consideration of the efficiency in actual operation, the upper limit of the heating time is preferably set to about 120 seconds.
[0051]
Further, the cooling rate after the heat treatment must be 5 ° C./s or more up to the hot-dip galvanizing treatment temperature from the viewpoint of promoting the martensitic transformation of the γ phase and suppressing the growth of crystal grains.
The heat treatment is desirably performed in a continuous galvanizing line.
[0052]
Subsequent to the above-mentioned heat treatment before plating, hot-dip galvanizing or alloying treatment is further performed to obtain a structure in which ferrite having an average crystal grain size of 3.5 μm or less is a main phase and has hard martensite as a second phase. High-strength hot-dip galvanized cold-rolled steel sheet can be obtained.
[0053]
In the hot-dip galvanizing treatment in the present invention, hot-dip galvanizing is performed in a temperature range of 450 to 550 ° C. in the same manner as in a normal hot-dip galvanizing line to form a hot-dip galvanized layer on the steel sheet surface.
Further, in the present invention, after the galvanizing treatment, an alloying treatment for alloying the galvanized layer can be performed. It is preferable that the processing temperature in this alloying process be approximately 470-570 ° C. If the heating temperature is lower than 470 ° C., the progress of alloying is slow and the productivity is lowered. On the other hand, if the heating temperature is higher than 570 ° C., the alloying of the coating layer proceeds too much and the galvannealed layer becomes brittle. Because.
[0054]
Furthermore, in order to secure a necessary amount of martensite after plating and to suppress grain growth, a cooling rate up to 300 ° C. after plating, or a cooling rate up to 300 ° C. after alloying when alloying is further performed. Needs to be 5 ° C./s or more.
Thus, by using the above manufacturing method, a hot-dip galvanized cold-rolled steel sheet having an ultrafine grain structure, an excellent balance between strength and ductility, and further excellent fatigue properties can be obtained.
[0055]
【Example】
The 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. In addition, the finish rolling temperature was set to _three or more Ar points in accordance with a conventional method. The hot-rolled sheet is pickled and then cold-rolled (rolling reduction: 60%) to form a 1.6-mm-thick cold-rolled sheet. Crystal annealing was performed, and then, in a continuous hot-dip galvanizing line, subsequent to heat treatment, plating treatment and alloying treatment were performed to obtain a product plate. At this time, the rate of temperature rise during the cold-rolled sheet annealing was 5 to 10 ° C./s. The bath temperature in the hot-dip galvanizing treatment was 465 ° C., and the alloying temperature in the case where the alloying treatment was performed thereafter was 520 ° C. In addition, No. No. 19 was only hot-dip galvanized and was not alloyed.
Table 3 shows the results obtained by examining the structure, tensile properties and fatigue properties of the product sheet thus obtained.
[0056]
The microstructure was observed by using an optical microscope or an electron microscope on the cross section in the rolling direction of the steel sheet, and the average crystal grain size of the ferrite was obtained, and the area ratio of each tissue was obtained, which was defined as the volume ratio. 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 collected from the rolling direction of the steel sheet.
Further, the fatigue characteristics were determined by a repetitive bending test of two swings using a test piece of JIS Z 2275 (a flat plate bending fatigue test method for a flat metal plate) having the dimensions and shapes shown in FIG. At this time, the point in time when the number of cycles reached 10 million was defined as the fatigue limit FL.
[0057]
[Table 1]

[0058]
[Table 2]

[0059]
[Table 3]

[0060]
As shown in Table 3, in each of the invention examples, the average grain size of the main phase, that is, ferrite occupying a fraction of 65% or more was as fine as 3.5 μm or less, and in particular, the amount of Ni and Mn was increased, and A ₃ No. using G steel with reduced Sample No. 16 is an ultrafine particle having an average crystal grain size of 0.9 μm. In addition, all of the invention examples are excellent in strength-ductility balance with TS × EL of 17500 MPa ·% or more, and the fatigue characteristics are such that the ratio between the fatigue limit FL and TS (FL / TS) is 0.45 or more. And a high level has been obtained. This is presumably because the existence of hard martensite as the second phase and the effect of grain refinement suppressed the occurrence of fatigue cracks and the propagation of cracks. Further, it is considered that the ferrite phase was hardened by fine precipitates such as TiC and NbC by setting the winding temperature within the above preferable range, but the fatigue characteristics could be further improved.
[0061]
On the other hand, no. In No. 5, the heating temperature before plating was low, and the austenite fraction during heating was low. As a result, the martensite fraction was low, and the TS × EL and FL / TS values were also small.
No. In No. 6, on the contrary, since the heating temperature before plating was too high, the concentration of the components to γ due to the two-phase separation did not progress during the heating, so that the martensite fraction was also low and the FL / TS value was lowered.
No. In No. 7, since the cooling rate before heating to the plating temperature after plating and the cooling rate after alloying were low, the second phase was mainly bainite softer than martensite, and the TS × EL and FL / TS values were inferior. Was.
No. In No. 11, since the heating temperature of the slab was low, TiC was coarsened, the effect of increasing the recrystallization temperature was suppressed, and the effect of reducing the crystal grain size of the steel sheet was not obtained, and the crystal grain size was increased. In addition, the TS × EL and FL / TS values are also reduced.
No. In No. 12, since the recrystallization annealing temperature greatly exceeded the proper upper limit temperature (846 ° C.) of the present invention, crystal grain growth was severe and TS × EL and FL / TS were deteriorated.
No. In Sample No. 13, the recrystallization was not completed because the annealing temperature was below the lower limit (816 ° C.) of the present invention, and the processed structure remained, so that the TS × EL value was extremely poor.
No. In No. 14, since the cooling rate after the recrystallization annealing was low, the crystal grains were coarsened, the strength was reduced, and the TS × EL value was deteriorated.
No. 23, T _X from less than A _1, gamma grain refining effect by recrystallization annealing can not be obtained, Kitaichi a decrease in strength due to a coarse grain, resulting TS × EL, the FL / TS value Led to a decline.
No. In No. 24, since A ₃ exceeded 860 ° C., high-temperature annealing was required, and as a result, crystal grains grew and the TS × EL value was reduced.
No. In No. 25, since the amount of (Ni + Mn) was small, the degree of supercooling during the γ-α transformation in the cooling process after annealing was small, and α could not be generated as fine nuclei, so that the crystal grains became coarse.
[0062]
【The invention's effect】
Thus, according to the present invention, a high-strength hot-dip galvanized cold-rolled steel sheet having an ultra-fine grain structure, having excellent mechanical properties in terms of strength-elongation balance, and also having excellent fatigue properties, has a large production facility. It can be manufactured stably without major remodeling, and is extremely useful in industry.
[Brief description of the drawings]
FIG. 1 is a diagram showing the relationship between the recrystallization temperature and the amounts of Ti and Nb added when the amounts of Ti and Nb are variously changed in a steel composition adjusted to A ₁ = 700 ° C. and A ₃ = 855 ° C. It is.
[Figure 2] ^{637.5 + 4930 {Ti * + (} 48/93) · [% Nb]} under the conditions of ≧ _{A 1,} the relationship between the recrystallization temperature Tre and _{A 3} with changes in _{A 3} to various FIG.
FIG. 3 is a view showing the dimensions and shape of a test piece for a plate bending fatigue test.

Claims (7)

In mass%,
C: 0.03 to 0.16%,
Si: 2.0% or less,
Mn: 3.0% or less and / or Ni: 3.0% or less,
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, and has a structure in which ferrite having an average crystal grain size of 3.5 μm or less is 65 vol% or more, and martensite is 10 vol% or more as a second phase, Furthermore, a hot-dip galvanized cold-rolled steel sheet having an ultra-fine grain structure and excellent fatigue properties, characterized by having a hot-dip galvanized layer on the surface.

[% M] is the content of element M (% by mass). 2. The steel sheet according to claim 1, further comprising:
Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. Cold rolled steel sheet. The steel sheet according to claim 1 or 2, further comprising:
Hot-dip galvanizing cold-rolling having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains a total of 0.005% or less of one or more selected from Ca, REM and B. steel sheet. In mass%,
C: 0.03 to 0.16%,
Si: 2.0% or less,
Mn: 3.0% or less and / or Ni: 3.0% or less,
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 obtained by heating a steel material having a composition of Fe and unavoidable impurities to 1200 ° C. or higher, performing hot rolling, winding, and then performing cold rolling, and then obtaining the following equation (6). The recrystallization annealing is performed at a temperature of not less than A ₃ (° C.) and not more than (A ₃ +30) (° C.), and then cooled to 600 ° C. at a rate of 5 ° C./s or more. Heat treatment is performed at a temperature in the range of (A ₁ +70) (° C.) or more and A ₃ (° C.) or less, and then cooled at a rate of 5 ° C./s or more to the hot-dip galvanizing treatment temperature to perform hot-dip galvanizing. After that, cool to 300 ° C at a rate of 5 ° C / s or more. Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that after the above-mentioned hot-dip galvanizing, the alloying treatment is further performed and then cooled to 300 ° C. at a rate of 5 ° C./s or more. Manufacturing method of cold rolled steel sheet.

[% M] is the content of element M (% by mass). The method for producing a hot-dip galvanized cold-rolled steel sheet according to claim 4, wherein the winding temperature is 750 ° C. or less and 650 ° C. or more, wherein the hot-dip galvanized steel sheet has an ultrafine grain structure and excellent fatigue properties. The steel material according to claim 4 or 5, further comprising:
Hot-dip galvanizing having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains one or two selected from Mo: 1.0% or less and Cr: 1.0% or less. Manufacturing method of cold rolled steel sheet. The steel material according to claim 4, 5 or 6, further comprising:
Hot-dip galvanizing cold-rolling having an ultrafine grain structure and excellent fatigue properties, characterized in that the composition contains a total of 0.005% or less of one or more selected from Ca, REM and B. Steel plate manufacturing method.

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