JP4107084B2 - Cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability and method for producing the same - Google Patents

Description

また、[%Ｍ] はＭ元素の含有量（質量％）
【００１８】
５．上記４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スポット溶接性に優れる冷延鋼板の製造方法。
【００１９】
６．上記４または５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スポット溶接性に優れる冷延鋼板の製造方法。
【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、セメンタイトを生成させて組織の高温安定性を高め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低下や溶接部の靱性低下を招くため、Ｃは0.03〜0.16％の範囲に限定した。
【００２１】
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。
【００２２】
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、Ａ₁ ，Ａ₃ 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。
【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。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。
【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。
【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。
【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。
【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。

また、[%Ｍ] はＭ元素の含有量（質量％）
【００３１】
なお、上記のＡ₁ , Ａ₃ はそれぞれ、鋼のＡc₁変態点温度（℃）、Ａc₃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。
【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Ａ₃ と等価になることが判明した。
【００３３】
図１に、Ａ₁ ＝700 ℃、Ａ₃ ＝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Ａ₁ 以上になると、再結晶温度Ｔreは 855℃近傍すなわちＡ₃ 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] ｝≧Ａ₁ の条件下において、Ａ₃ （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Ａ₃ 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] ｝≧Ａ₁ の条件下では、再結晶温度ＴreはＡ₃ 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、Ａ₁ 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、Ａ₃ 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Ａ₃ 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Ａ₃ 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。
【００３６】
次に、 (2)式は、Ａ₃ 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、Ａ₃ は実質的に再結晶温度になるため、Ａ₃ 以上の温度で再結晶焼鈍を行う必要がある。ここに、Ａ₃ 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、Ａ₃ ≦860 ℃を満足させる必要がある。なお、好ましくはＡ₃ ≦ 830℃である。
【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。
本発明では、鋼組織は、フェライト相の組織分率を体積率で85％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、セメンタイトを多量に析出させた組織にした上で、本発明で所期した強度、延性および強度−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85 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−伸びバランスが劣化し、また軟質なフェライトの組織分率が85 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。
【００３９】
また、フェライト以外の第２相組織としては、セメンタイトを0.3vol％以上 5.0 vol ％未満含有させることが重要である。
というのは、セメンタイトを微細に分散析出させることによって、スポット溶接時における熱影響部（ＨＡＺ部）のフェライト粒の粗大化、ひいては軟質化に伴う強度の低下を、効果的に抑制することができるからである。
しかしながら、セメンタイトの析出量が0.3vol％に満たないと、上記の効果が得難いので、セメンタイトは0.3vol％以上含有させることにした。但し、5.0vol％以上含有させようとすると、マルテンサイトなどの硬質第２相が出現し易くなるので、上限は5.0vol％程度が好適であり、 5.0vol ％未満とした。
【００４０】
なお、上記したフェライト相やセメンタイトの他、ベイナイト相やマルテンサイト相、パーライト相等が析出する場合もあるが、これらは合計量が10 vol％以下であれば、特に問題はない。
【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、熱間圧延を施し、ついで冷間圧延後、温度Ａ₃ （℃）以上、（Ａ₃ ＋30）（℃）以下で再結晶焼鈍を施し、その後 700℃まで５℃/s以上の速度で冷却し、ついで 650℃から 500℃までの冷却時間を30ｓ以上、 400ｓ以下とする。
【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、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、Ａr₃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、仕上げ圧延終了温度はＡr₃変態点以上とすることが好ましい。
【００４３】
熱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Ａ₃(℃）以上、（Ａ₃＋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、Ａ₃ が再結晶温度と等価となっているので、Ａ₃ 未満の温度では再結晶が不十分となる。一方、（Ａ₃ ＋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】
引き続き、焼鈍温度から 700℃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 700℃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 700℃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 700℃としたのは、結晶粒の微細化にはγ→α変態が開始する 700℃までが強く影響するからである。
【００４７】
さらに引き続き、 650℃から 500℃までの冷却時間（ 650℃から 500℃まで冷却するのに要する滞留時間) を30ｓ以上 400ｓ以下とする。上記の冷却時間が30ｓ未満では、第２相がベイナイトやマルテンサイトなどの硬質相となり易く、強度と延性は得られるものの、スポット溶接時の入熱によって容易に焼き戻されて、ＨＡＺ部が軟質化し、溶接強度が低下することとなる。また、冷却時間が 400ｓよりも長い場合は、結晶粒が粗大化するだけでなく、第２相が脆いパーライトになり易く、延性および伸びフランジ加工性が劣化する。
【００４８】
この点、650 ℃から 500℃までの冷却時間を30ｓ以上 400ｓ以下とすることにより、第２相に粒状のセメンタイトを析出させることができる。セメンタイトは高温まで比較的安定に存在するため、スポット溶接程度の小さな入熱であれば、ＨＡＺ部の粒成長を効果的に抑制して、ＨＡＺ軟化を防止することができる。
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延性およびスポット溶接性に優れる冷延鋼板を得ることができるのである。
【００４９】
【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、この際、巻取り温度は 570〜630 ℃とした。この熱延板を、酸洗後、冷間圧延（圧下率：60％）して、 1.6mm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、製品板とした。
かくして得られた製品板の組織、引張特性およびスポット溶接性について調査した結果を表３に示す。
【００５０】
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。
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JIS ５号試験片を用いた引張試験により測定した。
さらに、スポット溶接強度は、JIS Z 3136（スポット溶接継手の引張せん断試験方法）に準じて、溶接部のせん断強度（TSS)を測定した。
【００５１】
【表１】

【００５２】
【表２】

【００５３】
【表３】

【００５４】
表３に示したとおり、発明例はいずれも、85％以上の分率を占めるフェライトの平均粒径が 3.2μm 以下と微細であり、特にNi，Mn量を増量してＡ₃ を低下させたＧ鋼を用いたNo.14 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 17000 MPa・％以上と強度−延性バランスに優れ、さらにスポット溶接強度もJIS-Ａ級をはるかにしのいで高いレベルにあり、TSS/TS×100 の比も 3.0以上と高い数値であることが分かる。
【００５５】
これに対し、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値も小さくなっている。
No.10 は、焼鈍温度が本発明の適正温度（846 ℃）を大きく超えたため、結晶粒成長が激しく、TS×EL値が劣っている。
No.11 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値が劣っている。
No.12 は、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.21 は、Ｔ_X がＡ₁ 未満であるため、再結晶焼鈍によるα粒微細化効果が得られず、粗大粒となったため、十分な強度が得られなかった。
No.22 は、Ａ₃ が 860℃を超えたため、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.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、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。
No.3は、焼鈍後の 650℃から500 ℃までの冷却時間が不足したため、セメンタイトの析出が少なく、素材強度は高いものの、スポット溶接時の入熱でＨＡＺが軟質化し、TSS/TS×100 が 3.0未満まで低下した。
No.4は、焼鈍後の 650から500 ℃までの冷却時間が長すぎたため、結晶粒が粗大化すると共に、セメンタイトの析出が少なく、延性およびスポット溶接強度(TSS/TS ×100)が低下した。
【００５６】
【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、さらにはスポット溶接性に優れる高張力冷延鋼板を、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】 Ａ₁ ＝700 ℃、Ａ₃ ＝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] ｝≧Ａ₁ の条件下において、Ａ₃ を種々に変化させた場合におけるＡ₃ と再結晶温度Ｔreとの関係を示した図である。[0001]
BACKGROUND OF THE INVENTION
The present invention is a cold-rolled steel sheet suitable for use as an automobile, home appliance, and further machine structural steel, particularly a high-tensile cold-rolled steel sheet having an ultrafine grain structure and excellent in strength, ductility and spot weldability, and It relates to the manufacturing method.
[0002]
[Prior art]
Steel materials used as steel plates for automobiles, home appliances, and machine structures are required to have excellent mechanical properties such as strength and workability. As a means for comprehensively improving these mechanical properties, it is effective to make the structure finer, and so far, many manufacturing methods for obtaining a fine structure have been proposed.
[0003]
In recent years, with regard to high-strength steel, the goal is shifting to the development of high-strength steel sheets that can achieve both low cost and high-functional properties. Furthermore, steel plates for automobiles are required to have excellent impact resistance in addition to high strength in terms of protecting passengers in the event of a collision.
[0004]
Furthermore, since many automotive parts made of steel sheets are formed by press working, excellent formability is required as a steel sheet for automobile parts. In particular, parts used for members, reinforcements, etc., which are skeleton members for securing the strength of automobile bodies, are often molded with parts that use deformation of stretch flanges. A steel sheet for parts is also strongly required to have high stretch flangeability as well as high strength.
[0005]
Under such circumstances, miniaturization of the structure in high-strength steel is an important issue for the purpose of suppressing deterioration such as ductility and stretch flangeability associated with high tension.
[0006]
As a means for refining the structure, a large reduction rolling method is conventionally known. The main point of the structure refining mechanism in this large rolling reduction method is to apply a large rolling to the austenite grains to promote γ-α strain-induced transformation (see, for example, Patent Document 1).
Moreover, the case where the control rolling method and the control cooling method are applied is also known (for example, refer patent document 2).
[0007]
In addition, with regard to the raw steel, at least a part of the steel structure is made of ferrite, and the temperature is raised to a temperature range higher than the transformation point (Ac ₁ point) while plastic working is added thereto, or following this temperature rise, Ac Hold at a temperature range of _one or more points for a certain period of time, once part or all of the structure is reversely transformed into austenite, then ultrafine austenite grains appear, and then cooled to have an average crystal grain size of 5 μm or less, etc. There has been proposed a technique for forming a structure mainly composed of isotropic ferrite crystal grains (Patent Document 3).
[0008]
However, the driving force for crystal grain growth is proportional to the reciprocal of the grain size, so the crystal grains tend to grow more easily as the crystal grains become finer. However, when heat input such as spot welding is received in the subsequent assembly process, the crystal is coarsened in the heat affected (HAZ) part, which may cause a problem that joint strength and durability are lowered. .
In addition, even in the case of fine-grained steel, when the structure is strengthened by a hard second phase such as martensite or bainite, the second phase of the HAZ part softens due to heat input during welding, which also reduces the joint strength. Problem may occur.
[0009]
Since the conventional technology completely lacks consideration for these points, the high-strength steel sheet has good spot weldability that does not cause HAZ softening, and can exhibit the effect of fine graining even in the welded portion. Development was strongly desired.
[0010]
[Patent Document 1]
Japanese Patent Publication No. 5-65564 (Claims)
[Patent Document 2]
JP 63-128117 A (Claims)
[Patent Document 3]
JP-A-2-301540 (Claims)
[0011]
[Problems to be solved by the invention]
The present invention has been developed in view of the above-described situation, and it is possible to make ultra-fine grained cold-rolled steel sheets used as automobile, household appliance, and mechanical structural steel sheets, and only for strength and ductility. An object of the present invention is to propose a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, which effectively solves the problems related to spot weldability, together with its advantageous manufacturing method.
[0012]
Here, the target values of the strength-ductility balance and spot weldability of the cold-rolled steel sheet according to the present invention are as follows.
・ Strength-ductility balance (TS × El) ≧ 17000 MPa ・%
・ The shear strength (TSS) of the spot weld is more than the minimum value of JIS-A class, and TSS (kN) / TS (MPa) × 100 is more than 3.0. JIS-A class is JIS Z 3140 This is a grade applied to welds that require strength, and the minimum value of the shear load is determined by the tensile strength and thickness of the base metal.
[0013]
[Means for Solving the Problems]
As a result of intensive studies to solve the above problems, the inventors have adjusted the alloy elements appropriately to control the recrystallization temperature of the steel sheet and the A ₁ and A ₃ transformation temperatures, and then cold rolling. By optimizing the subsequent recrystallization annealing temperature and the subsequent cooling rate, the average crystal grain size of ferrite as the main phase becomes an ultrafine grain structure of 3.5 μm or less, and the second phase is optimized. The inventors have found that spot weldability (HAZ softening resistance) is significantly improved.
The present invention is based on the above findings.
[0014]
That is, the gist configuration of the present invention is as follows.
1. % By mass
C: 0.03-0.16%,
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 within the ranges satisfying the following formulas (1), (2) and (3) respectively, and the balance is becomes Fe and incidental impurities, further average crystal grain size has the following ferrite 3.5 [mu] m 85 vol% or more, and cementite possess less 0.3 vol% or more 5.0 vol% of a second phase, the ferrite and A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the total amount of phases other than cementite is 10 vol % or less .
Record
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] --- (4)
A ₁ = 727 + 14 [% Si] -28.4 [% Mn] -21.6 [% Ni] --- (5)
A ₃ = 920 + 612.8 [% C] ² − 507.7 [% C] + 9.8 [% Si] ³
−9.5 [% Si] ² +68.5 [% Si] +2 [% Mn] ² −38 [% Mn]
+ 2.8 [% Ni] ² − 38.6 [% Ni] +102 [% Ti] +51.7 [% Nb] --- (6)
[% M] is the content of M element (mass%)
[0015]
2. In 1 above, the cold-rolled steel sheet is further mass%,
Mo: 1.0% or less and
Cr: A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized by having a composition containing one or two selected from 1.0% or less.
[0016]
3. In the above 1 or 2, the cold-rolled steel sheet is further in mass%,
A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the composition contains one or more selected from Ca, REM and B in a total amount of 0.005% or less.
[0017]
4). % By mass
C: 0.03-0.16%,
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 within the ranges satisfying the following formulas (1), (2) and (3) respectively, and the balance is A steel material having a composition of Fe and inevitable impurities is heated to 1200 ° C. or higher, then hot-rolled, and then cold-rolled, and then a temperature A ₃ (° C.) or higher obtained by the following formula (6) (A ₃ +30) Recrystallization annealing at (° C) or lower, then cool to 700 ° C at a rate of 5 ° C / s or higher, and then set the cooling time from 650 ° C to 500 ° C to 30s or longer and 400s or shorter. A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability.

[% M] is the content of M element (mass%)
[0018]
5. In 4 above, the steel material is further mass%,
Mo: 1.0% or less and
Cr: A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, wherein the composition contains one or two selected from 1.0% or less.
[0019]
6). In the above 4 or 5, the steel material is further in mass%,
A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the composition contains one or more selected from Ca, REM and B in a total amount of 0.005% or less. .
[0020]
DETAILED DESCRIPTION OF THE INVENTION
The present invention will be specifically described below.
First, the reason why the composition of steel is limited to the above range in the present invention will be described. Unless otherwise specified, “%” in relation to ingredients means mass%.
C: 0.03-0.16%
C is not only an inexpensive reinforcing component, but also an element useful for generating cementite and improving the high temperature stability of the structure. However, if the content is less than 0.03%, the effect of addition is poor. On the other hand, if the content exceeds 0.16%, ductility and toughness of the welded portion are reduced, so C is limited to a range of 0.03 to 0.16%.
[0021]
Si: 2.0% or less
Si, as a solid solution strengthening component, effectively contributes to improving strength while improving the strength-elongation balance, but excessive addition deteriorates ductility, surface properties and weldability, so Si is 2.0% It was made to contain below. In addition, Preferably it is 0.01 to 0.6% of range.
[0022]
Mn: 3.0% or less and / or Ni: 3.0% or less
Both Mn and Ni are austenite stabilizing elements, contributing to the refinement of crystal grains through the action of lowering the A ₁ and A ₃ transformation points, and the strength-ductility balance through the action of promoting the formation of the second phase. Has the effect of increasing However, addition of a large amount hardens the steel and, on the other hand, deteriorates the strength-ductility balance.
Mn also has an effect of detoxifying harmful solid solution S as MnS, so it is preferably contained in an amount of 0.1% or more. Moreover, it is preferable to contain Ni 0.01% or more.
[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 the recrystallization temperature of the steel sheet is increased. For that purpose, it is preferable to contain each 0.01% or more. These may be added individually or in combination, but if they are added in excess of 0.2%, not only will the effect be saturated, but the amount of precipitates will increase and the ductility of the ferrite will increase. In any case, the content is 0.2% or less.
[0024]
Al: 0.01 to 0.1%
Al acts as a deoxidizer and is an element effective for the cleanliness of steel, and it is desirable to add it in the deoxidation process. Here, if the amount of Al is less than 0.01%, the effect of addition is poor. On the other hand, if it exceeds 0.1%, the effect is saturated, and rather the production cost is increased, so Al is limited to the range of 0.01 to 0.1%.
[0025]
P: 0.1% or less P is an element effective for achieving high strength at a low cost without causing a significant decrease in ductility. On the other hand, a large amount of P causes a decrease in workability and toughness. Was contained at 0.1% or less. In addition, when the request | requirement with respect to workability and toughness is severe, since it is preferable to reduce P, in this case, it is desirable to set it as 0.02% or less.
[0026]
S: 0.02% or less S is not only a cause of hot cracking during hot rolling, but also exists as an inclusion such as MnS in the steel sheet and causes deterioration of ductility and stretch flangeability, so it is desirable to reduce it as much as possible. However, up to 0.02% is acceptable, so in the present invention it was set to 0.02% or less.
[0027]
N: 0.005% or less Nitrogen causes aging deterioration and yield elongation, so it is suppressed to 0.005% or less.
[0028]
The basic components have been described above. 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 included as a reinforcing component as required. However, addition of a large amount deteriorates the strength-ductility balance on the contrary, so that it is desirable to contain each at 1.0% or less. In order to sufficiently exhibit the above-described action, it is preferable to contain 0.01% or more of Mo and Cr.
[0029]
0.005% or less total of one or more selected from Ca, REM and B
Ca, REM, and B all have the effect of improving workability through the control of the morphology of sulfides and the increase in grain boundary strength, and can be contained as required. However, excessive content may adversely affect cleanliness, so the total content is preferably 0.005% or less. In order to sufficiently exhibit the above-described action, it is preferable to contain 0.0005% or more of any one or more selected from Ca, REM, and B.
[0030]
As described above, the appropriate component composition range has been described. However, 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 contain the formulas (1), (2), and (3) in ranges that satisfy each.

[% M] is the content of M element (mass%)
[0031]
The above A ₁ and A ₃ are predicted values of the Ac ₁ transformation point temperature (° C.) and Ac ₃ transformation point temperature (° C.) of the steel, respectively, and are components derived from the detailed basic experiments of the inventors. It is a regression equation. This predicted temperature (° C.) is particularly suitable when applied at a heating rate of 2 ° C./s or more and 20 ° C./s or less.
[0032]
In the following, the reasons for limiting the above equations (1), (2), and (3) will be described in order.
Equation (1) is a condition that defines the amount of Ti and Nb added, and is based on the following findings.
In general, it is known that when Ti and Nb are added, TiC, NbC and the like are precipitated, and the recrystallization temperature of the steel sheet is increased. Therefore, a detailed investigation was made on the relationship between the addition amount of Ti and Nb and the recrystallization temperature Tre. When a certain amount or more of Ti and Nb was added, the recrystallization temperature became equivalent to A ₃ calculated by the above equation (6). It has been found.
[0033]
Fig. 1 shows the relationship between the amount of Ti and Nb added and the recrystallization temperature Tre when various amounts of Ti and Nb were added in the steel composition adjusted to A ₁ = 700 ° C and A ₃ = 855 ° C. Results are shown. Here, the recrystallization temperature Tre was determined by performing continuous annealing in a laboratory with various heating temperatures, measuring hardness, and observing the structure. The Ti addition amount uses Ti ^* as the effective Ti amount for precipitating TiC, and the Nb addition amount uses 48/93 · [% Nb] to convert it into Ti. The relationship with the crystal temperature is shown.
According to the figure, when 637.5 + 4930 {Ti ^* + (48/93) · [% Nb]} reaches 700 ° C, that is, A ₁ or more, the recrystallization temperature Tre rapidly rises to near 855 ° C, that is, near A ₃ and becomes saturated. I understand that.
[0034]
Next, in FIG. 2, A ₃ (varied by changing C, Si, Mn, Ni, etc.) under the condition of 637.5 + 4930 {Ti ^* + (48/93) · [% Nb]} ≧ A ₁ We are shown the results of examining the relationship between the recrystallization temperature Tre and a ₃ in the case of changing variously.
As shown in the figure, the recrystallization temperature Tre is equivalent to A ₃ under the condition of 637.5 + 4930 {Ti ^* + (48/93) · [% Nb]} ≧ A ₁ .
[0035]
Although this reason is not necessarily clear, it can be considered as follows.
That is, when Ti and Nb are added, the recrystallization temperature rises due to the pinning force of these fine carbides, and when recrystallization cannot be performed in the ferrite (α) region below A ₁ , the unrecrystallized processed α remains. (Ferrite + austenite (γ)) Competition between recrystallized α nucleation and α → γ transformation nucleation from processed α at preferential nucleation sites such as high dislocation density and non-uniform deformation at two-phase temperature range Occurs. At this time, since the driving force of the α → γ transformation is larger than the driving force of recrystallization, it is considered that γ nuclei are generated one after another in preference to the recrystallization α nucleation and occupy the preferential nucleation site.
Distortion (dislocation) is consumed by the atomic rearrangement in the α → γ transformation, and only the processed α having a low dislocation density remains, and recrystallization of the processed α becomes more difficult. The strain is completely eliminated only when the temperature rises, exceeds A _3, and enters the γ single phase region, and apparently recrystallization is completed. This is considered to be a mechanism in which the recrystallization temperature coincides with A ₃ and saturates.
Note that the α → γ transformation at this time nucleates from the processed α (many preferential nucleation sites), so that the γ grains at a high temperature at which recrystallization is completed are refined. Therefore, it is extremely effective to adjust the recrystallization temperature to A ₃ for refining high-temperature γ grains during annealing. Therefore, in the present invention, Ti and Nb satisfying the formula (1) are added. is there.
[0036]
Next, equation (2) is a condition for defining A ₃ .
As described above, when the expression (1) is satisfied, A ₃ substantially reaches the recrystallization temperature. Therefore, it is necessary to perform recrystallization annealing at a temperature equal to or higher than A ₃ . Here, when A ₃ exceeds 860 ° C., it is necessary to apply a recrystallization annealing temperature at a higher temperature, and γ grain growth is intense, and as a result, fine grains having an average crystal grain size of 3.5 μm or less were not obtained. . Therefore, it is necessary to satisfy A ₃ ≦ 860 ° C. Preferably, A ₃ ≦ 830 ° C.
[0037]
Next, equation (3) is a condition that defines the amount of Mn or Ni, that is, the amount of austenite stabilizing element added.
As the austenite stabilizing element increases, the ferrite start line in the CCT diagram shifts to the low temperature side, which increases the degree of subcooling during the γ → α transformation in the cooling process after annealing, and α nucleates finely. As a result, the α crystal grains are refined. Here, in order to obtain fine grains with an average crystal grain size of 3.5 μm or less, [% Mn] + [% Ni] ≧ 1.3 (%) in addition to the above formulas (1) and (2) There was a need.
In addition, as long as [% Mn] + [% Ni] ≧ 1.3 (%) is satisfied, Mn and Ni may be added alone or in combination. More preferably, the range is [% Mn] + [% Ni] ≧ 2.0 (%).
[0038]
Next, the steel structure will be described.
In the present invention, the steel structure has a ferrite phase with a volume fraction of 85% or more and an average crystal grain size of ferrite of 3.5 μm or less.
This is because, after making a structure in which a large amount of cementite is precipitated, in order to obtain a cold-rolled steel sheet excellent in strength, ductility and strength-elongation balance as expected in the present invention, a steel structure mainly composed of fine ferrites is used. This is because, in particular, it is important that the structure fraction of the fine ferrite phase having an average crystal grain size of 3.5 μm or less is 85 vol% or more.
Here, if the average crystal grain size of ferrite exceeds 3.5 μm, the strength-elongation balance deteriorates, and if the structure fraction of soft ferrite is less than 85 vol%, the ductility is significantly reduced and the workability becomes poor. .
[0039]
In addition, as the second phase structure other than ferrite, it is important to contain cementite in an amount of 0.3 vol % or more and less than 5.0 vol % .
This is because, by finely dispersing and precipitating cementite, it is possible to effectively suppress the coarsening of ferrite grains in the heat-affected zone (HAZ zone) during spot welding, and consequently the decrease in strength accompanying softening. Because.
However, if the precipitation amount of cementite is less than 0.3 vol%, it is difficult to obtain the above effect, so it was decided to contain cementite in an amount of 0.3 vol% or more. However, if an attempt is contained more than 5.0 vol%, since the hard second phase such as martensite is liable to appear, the upper limit is about 5.0 vol% is Ri preferred der was less than 5.0 vol%.
[0040]
In addition to the ferrite phase and cementite described above, a bainite phase, a martensite phase, a pearlite phase, and the like may be precipitated, but there is no particular problem if the total amount is 10 vol% or less.
[0041]
Next, manufacturing conditions will be described.
The steel adjusted to the above-mentioned preferred component 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 in a high temperature state or once cooled and then heated to 1200 ° C. or higher, and then hot-rolled, and after cold rolling, temperature A ₃ (° C.) or higher, (A ₃ +30) Recrystallization annealing is performed at (° C) or less, and then cooling is performed at a rate of 5 ° C / s or more to 700 ° C, and then the cooling time from 650 ° C to 500 ° C is set to 30 s or more and 400 s or less.
[0042]
In the above process, when the heating temperature of the slab is less than 1200 ° C., TiC and the like are not sufficiently dissolved and coarsened, and the effect of increasing the recrystallization temperature and the effect of suppressing grain growth in the subsequent recrystallization annealing process are insufficient. Therefore, the heating temperature of the slab needs to be 1200 ° C. or higher.
In the present invention, the hot finish rolling outlet temperature is not particularly limited, but if it is less than the Ar ₃ transformation point, α and γ are generated during rolling, and a band-like structure is likely to be formed in the steel sheet. Such a band-like structure remains even after cold rolling or annealing, and may cause anisotropy in material properties. Therefore, it is preferable that the finish rolling end temperature is equal to or higher than the Ar ₃ transformation point.
[0043]
The coiling temperature after completion of hot rolling is not particularly limited, but if it is less than 500 ° C or more than 650 ° C, the precipitation of AlN to suppress aging deterioration due to nitrogen is insufficient and the material properties are inferior. It becomes. Further, in order to make the structure of the steel sheet uniform and make the crystal grain size as fine and uniform as possible, the coiling temperature is preferably 500 ° C. or higher and 650 ° C. or lower.
[0044]
Next, preferably after removing the oxidized scale on the surface of the hot-rolled steel sheet by pickling, it is subjected to cold rolling to obtain a cold-rolled steel sheet having a predetermined thickness. Here, pickling conditions and cold rolling conditions are not particularly limited, and may be according to 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. On the other hand, if the rolling reduction is increased too much, Since 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 A ₃ (° C.) or more and (A ₃ +30) (° C.) or less shown in the above equation (6), and recrystallization annealing is performed.
In the steel material of the present invention whose components are adjusted as described above, since A ₃ is equivalent to the recrystallization temperature, recrystallization becomes insufficient at a temperature lower than A ₃ . On the other hand, at a temperature exceeding (A ₃ +30) (° C.), the growth of γ grains during annealing is intense and 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 a time shorter than 10 seconds, and a processed structure that remains stretched in the rolling direction and a recovered structure that has not been recrystallized remain, so that sufficient ductility may not be ensured. On the other hand, if the time is longer than 120 seconds, the γ crystal grains are coarsened and the desired strength may not be obtained.
[0046]
Subsequently, cooling is performed from the annealing temperature to 700 ° C. under a cooling rate of 5 ° C./s or more. Here, the cooling rate is an average cooling rate from the annealing temperature to 700 ° C. Here, when 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 700 ° C. needs to be 5 ° C./s or more.
The reason why the end point temperature of the above controlled cooling treatment is set to 700 ° C. is that the refinement of crystal grains is strongly influenced by 700 ° C. at which the γ → α transformation starts.
[0047]
Furthermore, the cooling time from 650 ° C. to 500 ° C. (the residence time required for cooling from 650 ° C. to 500 ° C.) is set to 30 s or more and 400 s or less. If the cooling time is less than 30 s, the second phase tends to be a hard phase such as bainite and martensite, and although strength and ductility are obtained, it is easily tempered by heat input during spot welding, and the HAZ part is soft. As a result, the welding strength is reduced. In addition, when the cooling time is longer than 400 s, not only the crystal grains become coarse, but the second phase easily becomes brittle pearlite, and ductility and stretch flangeability deteriorate.
[0048]
In this respect, by setting the cooling time from 650 ° C. to 500 ° C. to 30 s or more and 400 s or less, granular cementite can be precipitated in the second phase. Since cementite exists relatively stably up to a high temperature, if the heat input is as small as spot welding, grain growth in the HAZ part can be effectively suppressed and HAZ softening can be prevented.
Thus, by using the above manufacturing method, a cold-rolled steel sheet having an ultrafine grain structure and excellent in ductility and spot weldability can be obtained.
[0049]
【Example】
A slab having the component composition shown in Table 1 was heated under the conditions shown in Table 2 and hot-rolled according to a conventional method to obtain a 4.0 mm thick hot-rolled sheet. At this time, the winding temperature was 570 to 630 ° C. This hot-rolled sheet is pickled and cold-rolled (rolling ratio: 60%) to form a 1.6 mm-thick cold-rolled sheet, and then recrystallized and annealed under the conditions shown in Table 2 in a continuous annealing line. To make a product plate.
Table 3 shows the results of investigation on the structure, tensile properties and spot weldability of the product plate thus obtained.
[0050]
In addition, the structure was observed using an optical microscope or an electron microscope for the cross section in the rolling direction of the steel sheet, and the average crystal grain size of ferrite was determined, and the area ratio of each organization was determined and used as the volume ratio. Here, the average crystal grain size of ferrite was determined in accordance with the cutting method specified in JIS G 0552.
Tensile properties (tensile strength TS, elongation EL) were measured by a tensile test using JIS No. 5 test pieces collected from the rolling direction of the steel sheet.
Furthermore, the spot weld strength was determined by measuring the shear strength (TSS) of the weld according to JIS Z 3136 (Tensile shear test method for spot welded joints).
[0051]
[Table 1]

[0052]
[Table 2]

[0053]
[Table 3]

[0054]
As shown in Table 3, in all of the inventive examples, the average particle diameter of ferrite occupying a fraction of 85% or more is as fine as 3.2 μm or less, and in particular, the amount of Ni and Mn was increased to decrease A ₃ . No.14 using G steel has an average grain size of 0.9μm and is very fine. In all of the invention examples, TS × EL is 17000 MPa ·% or more and the strength-ductility balance is excellent, and the spot welding strength is far higher than JIS-A class, and TSS / TS × 100 It can be seen that the ratio is also a high value of 3.0 or more.
[0055]
On the other hand, in No. 9, since the heating temperature of the slab was low, TiC was coarsened, the effect of increasing the recrystallization temperature was suppressed, and the grain refinement effect of the steel sheet was not obtained, and the crystal grain size was increased. . Also, the TS × EL value is small.
In No. 10, since the annealing temperature greatly exceeded the appropriate temperature (846 ° C.) of the present invention, the crystal grain growth was intense and the TS × EL value was inferior.
In No. 11, since the annealing temperature was less than the lower limit (816 ° C.) of the present invention, the recrystallization was not completed and the processed structure remained, so the TS × EL value was inferior.
In No. 12, since the cooling rate after annealing was low, the crystal grains became coarse and the strength decreased, leading to deterioration of the TS × EL value.
In No. 21, since T _X was less than A ₁ , the α grain refinement effect by recrystallization annealing was not obtained, and coarse grains were obtained, so that sufficient strength was not obtained.
In No. 22, since A ₃ exceeded 860 ° C., high temperature annealing was required. As a result, crystal grains grew and the TS × EL value decreased.
In No. 23, since the amount of (Ni + Mn) is small, the degree of supercooling during the γ-α transformation in the cooling process after annealing was small, and α could not produce fine nuclei, so the crystal grains became coarse .
In No.3, the cooling time from 650 ° C to 500 ° C after annealing was insufficient, so there was little cementite precipitation and the material strength was high, but the HAZ softened by heat input during spot welding, and TSS / TS × 100 Decreased to less than 3.0.
In No.4, since the cooling time from 650 to 500 ℃ after annealing was too long, the crystal grains became coarse, the precipitation of cementite was small, and the ductility and spot weld strength (TSS / TS × 100) decreased. .
[0056]
【The invention's effect】
Thus, according to the present invention, a high-tensile cold-rolled steel sheet having an ultra-fine grain structure, excellent mechanical strength and excellent strength-elongation balance, and further excellent spot weldability is accompanied by a significant remodeling of manufacturing equipment. It can be stably produced without any trouble and is extremely useful in industry.
[Brief description of the drawings]
FIG. 1 is a graph showing the relationship between Ti and Nb addition amounts and recrystallization temperature when Ti and Nb addition amounts are variously changed in a steel composition adjusted to A ₁ = 700 ° C. and A ₃ = 855 ° C. It is.
In Figure 2 ^{637.5 + 4930 {Ti * + (} 48/93) · [% Nb]} conditions ≧ A _1, the relationship between the recrystallization temperature Tre and A ₃ with changes in A ₃ to various FIG.

Claims (6)

% By mass
C: 0.03-0.16%,
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 within the ranges satisfying the following formulas (1), (2) and (3) respectively, and the balance is becomes Fe and incidental impurities, further average crystal grain size has the following ferrite 3.5 [mu] m 85 vol% or more, and cementite possess less 0.3 vol% or more 5.0 vol% of a second phase, the ferrite and A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the total amount of phases other than cementite is 10 vol % or less .
Record
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] --- (4)
A ₁ = 727 + 14 [% Si] -28.4 [% Mn] -21.6 [% Ni] --- (5)
A ₃ = 920 + 612.8 [% C] ² − 507.7 [% C] + 9.8 [% Si] ³
−9.5 [% Si] ² +68.5 [% Si] +2 [% Mn] ² −38 [% Mn]
+ 2.8 [% Ni] ² − 38.6 [% Ni] +102 [% Ti] +51.7 [% Nb] --- (6)
[% M] is the content of M element (mass%) In claim 1, the cold-rolled steel sheet is further in mass%,
Mo: 1.0% or less and
Cr: A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized by having a composition containing one or two selected from 1.0% or less. The cold-rolled steel sheet according to claim 1 or 2, further in mass%,
A cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the composition contains one or more selected from Ca, REM and B in a total amount of 0.005% or less. % By mass
C: 0.03-0.16%,
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 within the ranges satisfying the following formulas (1), (2) and (3) respectively, and the balance is A steel material having a composition of Fe and inevitable impurities is heated to 1200 ° C. or higher, then hot-rolled, and then cold-rolled, and then a temperature A ₃ (° C.) or higher obtained by the following formula (6) (A ₃ +30) Recrystallization annealing at (° C) or lower, then cool to 700 ° C at a rate of 5 ° C / s or higher, and then set the cooling time from 650 ° C to 500 ° C to 30s or more and 400s or less. A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability.

[% M] is the content of M element (mass%) In claim 4, the steel material is further in mass%,
Mo: 1.0% or less and
Cr: A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, wherein the composition contains one or two selected from 1.0% or less. In Claim 4 or 5, steel material is further mass%,
A method for producing a cold-rolled steel sheet having an ultrafine grain structure and excellent spot weldability, characterized in that the composition contains one or more selected from Ca, REM and B in a total amount of 0.005% or less. .

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