JP2004100020A - Strain age hardening type hot rolled steel structural member having excellent impact property and method for producing the hot rolled steel member - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a strain age hardening type hot rolled steel structural member which has satisfactory workability and excellent impact properties, and to provide a method for producing the hot rolled steel member. <P>SOLUTION: The structural member is obtained by subjecting a hot rolled steel member to forming, and electrodeposition coating/baking is to be applied thereto. The structural member has a composition comprising, by mass, 0.0005 to 0.2% C, 0.001 to 2.0% Si, 0.001 to 2.0% Mn, ≤0.2% P, ≤0.015% S, 0.002 to 2.0% Al, 0.0005 to 0.2% N and 0.001 to 2.0% Cu, and in which the Cu content is ≥0.l times the total of the C content and N content, and the balance Fe with inevitable impurities. When heat treatment equivalent to electrodeposition coating/baking at 170°C for 20 min is performed to the structural member, the variation in yield strength (YS<SB>-</SB>BHO) before and after the heat treatment in the part which is not deformed by the forming is ≥30 MPa. <P>COPYRIGHT: (C)2004,JPO

Description

【０００１】
【発明の属する技術分野】
本発明は、自動車の構造部材・足廻り部材・パネル部材の使途に好適である衝撃特性および成型加工性に優れた歪時効硬化型熱延鋼構造部材およびその熱延鋼材を製造する方法に関するものであり、引張強度で２５０ＭＰａから９００ＭＰａ程度の幅広い強度の鋼材に適用が可能である。本発明はさらに優れた疲労特性も併せ持つことから、上記用途の他にタイヤホイール用途、建築物・船舶・橋梁等の構造物用途にも適用することが可能である。
【０００２】
【従来の技術】
地球環境保護の観点から自動車が排出するＣＯ_２軽減が重要な課題となっている。ＣＯ_２軽減のためには車体重量の軽減が有効であり、そのために鋼材高強度化のニーズが高まっている。ところが、一般的に材料の高強度化は形状凍結性の低下や成型時の割れといったプレス成形性の劣化を伴うことが知られており、加工性を低下させずに高強度化する方法が強く望まれていた。自動車用鋼材に求められる他の特性としては、走行時の振動に対し材料劣化を起こさない特性、すなわち疲労特性や、高速衝突時において材料側でエネルギーを吸収し乗員の安全性を確保するための特性、すなわち耐衝撃性といった特性が挙げられる。
【０００３】
建築物や橋梁や船舶に用いられる構造用鋼材についても、構造物の軽量化を目的とした高強度鋼材の要請が高い。前記の自動車用鋼材と同様に、成形加工性や構造物同士の衝突を前提とした耐衝撃性、疲労特性が強く望まれている。
【０００４】
このような要望に対し、成形加工性を確保した上で高強度化を達成する技術として、成型加工時には軟質に保たれ、成形加工後の１７０℃×２０分程度の電着塗装焼付工程でおこる歪時効硬化現象を利用して降伏強度あるいは引張強度を増加させる、いわゆる焼付硬化性（Ｂａｋｅ　Ｈａｒｄｅｎａｂｉｌｉｔｙ：ＢＨ）を利用した技術が知られている。この種の鋼材は、成形加工時には炭素原子あるいは窒素原子を固溶させて成形性を確保しておき、電着塗装焼付工程において成形加工時に鋼材内に生じた転位に炭素原子あるいは窒素原子を固着させるか、あるいは転位上に炭化物あるいは窒化物を微細分散析出させることによって、降伏強度あるいは引張強度の上昇を図るものである。
【０００５】
一方、衝撃特性に関しては、高歪み速度での変形における変形応力が高いほど、衝撃変形時により多くのエネルギーが吸収されると考えられており、歪み時効硬化現象を利用して耐衝撃性を向上させる技術が提案されている。例えば、下記特許文献１には、所定の鋼成分において、固溶状態のＣおよびＮ量とフェライト結晶粒径を制御することにより降伏応力と引張強度の双方を上昇させた耐衝撃性に優れた歪時効硬化型熱延鋼板およびその製造方法が提案されている。また、下記特許文献２には、フェライト粒径を制御した、固溶Ｎ利用型の耐衝撃性に優れた歪時効硬化型高張力熱延鋼板およびその製造方法が提案されている。
【０００６】
【特許文献１】
特開２００１−３３５８８９号公報
【０００７】
【特許文献２】
特開２００１−２２６７４４号公報
【０００８】
しかし、部材の加工は曲げ加工、プレス加工などにより成形加工されることが多いが、実際の部材では未成形加工部分（未変形部分）が部材体積のかなりの割合を占め、そのため、（１）　変形を受けた部分をいかに歪み硬化現象で強化しても、結局はこの軟弱な未変形部分の強度で耐衝撃特性が決まってしまう、また（２）　これまでに提案されている歪み時効硬化型鋼材を成形加工→電着塗装焼付しても衝撃変形が加わった時の衝撃破壊の仕方に再現性が無く、このため実部材への適用が実質的に難しかったという問題点を有していた。
【０００９】
【発明が解決しようとする課題】
本発明は上記の如き実状に鑑みてなされたものであって、電着塗装工程を経て作られる自動車用の構造部材・足廻り部材・パネル部材用途、建築用の構造部材、電機製品の内外板パネルに好適な、良好な加工性を有しかつ衝撃特性に優れた歪時効硬化型熱延鋼構造部材およびその熱延鋼材の製造方法を提供することを目的とする。
【００１０】
【課題を解決するための手段】
本発明者らは、前記課題を解決するため、実部材の高速圧縮変形試験を行った結果、衝撃圧壊変形時に不安定変形を起こす原因は成形加工→電着塗装焼付工程後において部材内の強度の不均一が大きいことであり、さらに優れた耐衝撃性が得られにくい原因は、変形部分と未変形部分の変形応力の差が大きくなると、高速変形時の吸収エネルギーが結局未変形部分の特性で決まってしまうためであると考えるに至った。さらに電着塗装焼付処理工程を行った後の構造部材において大きな強度不均一が生じる原因としては、成形加工により局所的に変形が加わった部分では転位が導入されて歪時効硬化する一方で、未変形の部分では転位が導入されないため焼付塗装工程でも時効硬化が起こらないか、あるいはその硬化量がわずかであることが主因であるとの結論に到った。
【００１１】
そこで本発明者らは、実構造部材に成形加工したときに優れた耐衝撃特性を有する歪み時効硬化型熱延鋼材を達成すべく、鋭意実験と検討を重ねた結果、電着焼付塗装工程後において未変形部分でも所定範囲内の強度上昇を得ること、さらに電着焼付塗装工程後の変形部分と未変形部分の強度差を小さくすること、すなわち強度不均一を極力小さくすることが耐衝撃特性向上に有効であることを見出した。
そして、本発明者らは変形部分と未変形部分の強度差を小さくするためには、未変形部分のフェライト粒内でも炭素あるいは窒素原子のクラスタリングおよび析出を起こさせることが有効であることを見出し、またそのような組織にするためには、鋼中のＣｕ／（Ｃ＋Ｎ）成分比を適正値以上にすることが有効であることを明らかにした。
【００１２】
このような現象が発現するメカニズムを詳細に解析した結果、Ｃｕを適正量添加することによりＣｕ原子あるいはＣｕ添加により導入される空孔集合体が核生成サイトになり、転位が存在していない場所でも電着塗装焼付処理中に微細な炭窒化物あるいは炭窒素クラスターが析出するという全く新しい事実を見出した。
【００１３】
さらに、本発明者らは塗装焼付処理中に粒内に新たに再析出する微細な炭窒化物、炭窒素クラスターのＣおよびＮ供給源を調査した結果、（１）粒内に固溶しているＣあるいはＮ、（２）粒界に偏析しているＣあるいはＮといった過去に提案されている供給源の他に、（３）焼付塗装前のフェライト粒内に存在している不安定化した炭窒化物が供給源になりうるという全く新しい事実を初めて明らかにした。具体的には、析出物を成形加工前の段階で熱間圧延工程、冷間加工工程、調質圧延工程あるいはレベラー加工等によりエネルギー的に不安定化させておけば塗装焼付工程で時効硬化に寄与するＣおよびＮの供給源となることを初めて見出した。
【００１４】
本発明者らは、上記のように加工転位が存在しなくても炭素・窒素集合体あるいは炭窒化物を微細に析出させる効果のあるＣｕの適正量添加技術と、炭素・窒素集合体あるいは炭窒化物への炭素、窒素の供給源として不安定析出物を利用する技術を相互に利用することにより、未変形領域でも電着塗装相当の熱処理で十分な硬化をし、かつ成形加工による変形部分と未変形部分の強度差を小さくすることが可能であることを見出し、本発明をなすに到った。
【００１５】
また、本発明者らの検討の結果、上記した組成とすることにより、フェライト単相の組織の他に、フェライトを主体とする組織にマルテンサイト、残留オーステナイト、パーライトのうち１種または２種以上の組織を含む複相組織鋼においても、衝撃特性および成型加工性に優れた歪時効硬化型熱延鋼材を得られ、また強度レベルでいうと引張強度で２５０〜９００ＭＰａの広い範囲内で、従来鋼に比較して優れた衝撃特性が得られることがわかった。
【００１６】
本発明は、前記課題を解決するために次の手段を講じた。すなわち、
（１）第１の発明は、熱延鋼材を成形加工した電着塗装焼付処理を行う前の構造部材であって、質量％で、
Ｃ　：０．０００５〜０．２％、　Ｓｉ：０．００１〜２．０％、
Ｍｎ：０．００１〜２．０％、　　Ｐ　：０．２％以下、
Ｓ　：０．０１５％以下、　　　　Ａｌ：０．００２〜２．０％、
Ｎ　：０．０００５〜０．２％、　　Ｃｕ：０．００１〜２．０％
を含み、かつＣｕ含有量がＣ含有量とＮ含有量の和の０．１倍以上であり、残部がＦｅおよび不可避的不純物からなり、１７０℃×２０分間の電着塗装焼付相当の熱処理を行った際に、成形加工により変形していない部分における熱処理前後の降伏強度変化量（ＹＳ＿ＢＨ０）が３０ＭＰａ以上であることを特徴とする。
【００１７】
（２）第２の発明は、フェライト面積率が５０％以上であることを特徴とする前記（１）記載の歪時効硬化型熱延鋼構造部材であり、
（３）第３の発明は、前記（１）又は（２）記載の発明に加え、前記熱処理を行った際に、成形加工により変形した部分と、変形していない部分の降伏強度ＹＳ及び引張強度ＴＳの差が、それぞれＹＳで９０ＭＰａ以下かつＴＳで５０ＭＰａ以下であることを特徴とする。
【００１８】
（４）第４の発明は、前記（１）〜（３）の何れか１項に記載の発明に加え、前記熱処理を行った際に、未変形部フェライト粒内における炭素・窒素原子集合体および炭窒化物析出物の２種合計の分布密度が、体積密度で１００〜１００００個・μｍ^−３であることを特徴とし、
（５）第５の発明は、前記（１）〜（４）の何れかに記載の組成に加え、質量％で、Ｎｉ：０．１〜２．０％を含むことを特徴とする。
【００１９】
（６）第６の発明は、前記組成に加えてさらに、質量％で、下記ａ群〜ｄ群の１群または２群以上を含むことを特徴とする。
ａ群：Ｃｒ、Ｍｏ、Ｗのうち１種または２種以上の合計を０．１〜２．０％
ｂ群：Ｎｂ、Ｔｉ、Ｖ、Ｔａのうち１種または２種以上の合計を０．０３〜　０．２％
ｃ群：Ｂを０．０００３〜０．０１０％
ｄ群：ＣａまたはＭｇのうち１種または２種を合計で０．００１〜０．０１％
（７）第７の発明は、前記（１）〜（６）のいずれかに記載の熱延鋼材に電気めっき又は溶融めっきが施されていることを特徴とする衝撃特性に優れた歪時効硬化型熱延めっき鋼材である。
【００２０】
更に本発明は、衝撃特性に優れた歪時効硬化型熱延鋼材の製造方法であって、
（８）第８の発明は、前記（１）、（５）又は（６）記載の組成からなる鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせによりＡ１　温度〜１００℃間を平均冷却速度０．０１〜３０℃／ｓで冷却することを特徴とし、
（９）第９の発明は、（１）、（５）又は（６）記載の組成からなる鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせによる冷却の後、Ａ１　温度〜１００℃間で巻取ることを特徴とする。
【００２１】
（１０）第１０の発明は、前記（１）、（５）又は（６）記載の組成からなる鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせにより室温まで冷却したのち、１００℃〜Ａ１　温度の間で再加熱処理を行うことを特徴とし、
（１１）第１１の発明は、前記（８）〜（１０）の何れかに示した製造工程に続いて、伸び率：０．５〜４０％の調質圧延またはレベラー加工を施すことを特徴とする。
【００２２】
実際の成形加工においては、予歪量やその変形モード（引張か圧縮か、あるいは１軸引張か２軸引張か等）は部材の場所ごとに異なるが、本発明者らが様々な成形加工形式と歪み時効硬化挙動と衝撃特性との関係を詳細に検討した結果、２％の引張歪みで変形部の特性が良く代表できることを突き止めた。この知見を元に、本発明では、変形部分の予変形量として引張歪２％に定めて評価するものとする。なお、未変形部分とは素材を成形加工し、構造部材とする際の付加歪みがゼロであることを意味し、鋼材素材の製造中に付加されるロール圧延やダイス圧延による変形が加わっていても構わない。降伏強度については、耐衝撃特性と相関がある下降伏強度を指すものとする。ＹＳ＿ＢＨ０値、ＹＳ＿ΔＢＨ値、ＴＳ＿ΔＢＨ値の定義の模式図を図１に示す。
【００２３】
焼付塗装条件としては、実施例においては標準で用いられている１７０℃×２０分の熱処理でその評価を行った。ただし、本発明が有効な成形加工後の焼付塗装条件は、１７０℃×２０分に限定されるべきものではなく、ＹＳ＿ＢＨ０として３０ＭＰａ以上が得られるのであれば、焼付温度が１７０℃以下、例えば１００℃の低温焼付条件でも、また焼付時間も２０分以下の短時間保持でも本発明の目的を達成することができる。
【００２４】
本発明において「炭素・窒素原子集合体」とは、炭素あるいは窒素から構成される原子集合体であり、炭素原子と窒素原子を合わせて５個以上の原子からなるものを指す。また「フェライト」とは、下記の非特許文献１に示すようなポリゴナルフェライト組織、擬ポリゴナルフェライト組織、あるいはＭ／Ａ複合体を含むグラニュラーベイニティックフェライト組織を指す。
【００２５】
【非特許文献１】
ＩＳＩＪ　ｉｎｔｅｒｎａｔｉｏｎａｌ、３５巻（２００２），９４１〜　９４４頁
【００２６】
鋼中のＣｕの分析方法については、０．００２％を超える濃度の場合はスパーク放電発光分光分析法を用いることが簡易で好適であるが、０．００２％以下の場合には、ＩＣＰ発光分析法あるいは２次イオン質量分析法を用いる必要がある。なお、ＩＣＰ発光分析法を用いる場合には、ｓｏｌ．Ｃｕとｉｎｓｏｌ．Ｃｕを併せた全Ｃｕ量を測定する必要がある。
【００２７】
上記した本発明の熱延鋼材は各種めっき用原材として好適である。めっき層の形成は電気めっき法、溶融めっき法のいずれでも良く、めっきの主成分としては亜鉛、クロム、錫、ニッケルが例として挙げられる。
また本発明鋼は、優れた耐衝撃性の他に構造部材として成形加工したときに優れた耐疲労特性を有する特徴を併せ持つ。
【００２８】
【発明の実施の形態】
以下に、本発明について詳細に説明する。
まず成分の限定理由について説明する。成分含有量は質量％である。
Ｃ：Ｃは鋼の歪み時効硬化の発現およびミクロ組織の制御に必須の添加元素である。しかし、０．０００５％未満であると未変形部分における塗装焼付処理前後の降伏強度の変化量（ＹＳ＿ＢＨ０）＝３０ＭＰａ以上が達成できない。また、０．２％を超えると耐衝撃特性が急激に悪化し、また溶接性も低下する。このため本発明ではＣの範囲を０．０００５〜０．２％に限定した。
【００２９】
Ｓｉ：Ｓｉは鋼材のミクロ組織および強度の調整に用いられるので、０．００１％以上含有するものとする。しかしながら、２．０ｗｔ％を超えると熱延時の脱スケール性の悪化や、疲労特性の悪化を招く。従ってＳｉ含有量を２．０％以下の範囲に制限した。
【００３０】
Ｍｎ：Ｍｎは鋼材のミクロ組織および強度の調整に用いられるので、０．００１％以上含有するものとする。しかしながら、２．０％を超えると成形加工性の劣化を招く。従って、Ｍｎ含有量を２．０％以下の範囲に制限した。Ｓの熱間脆性を抑制させる意味では、０．００５％以上の添加が望ましい。
【００３１】
Ｐ：不純物であるＰは鋼材の強度の調整に用いられる。しかしながら、０．２％を超えると成形加工時の割れを起こす可能性があるので、Ｐ含有量の範囲を０．２％以下とした。
【００３２】
Ｓ：不純物であるＳはＭｎＳ、ＣｕＳとして鋼中に存在させ、結晶粒径の制御を通じて鋼材の強度・延性の調整に用いられる。また粒内に微細に分布させることで塗装焼付熱処理中に鋼中に再析出するＣおよびＮの供給源としても作用する。しかしながら、０．０１５％を超えると熱間脆性を起こす可能性があるので、その範囲を０．０１５％以下に限定した。
【００３３】
Ａｌ：Ａｌの重要な作用として、塗装焼付処理前にフェライト粒内にＡｌ（ＮＣ）として析出していると、塗装焼付処理時に再析出する炭素・窒素集合体あるいは炭窒化物形成のＮ、Ｃ供給源として作用し、その結果Ｃｕと同時添加することで未変形領域のＢＨが大きくなる。しかしながら、Ａｌの添加量が０．００２％未満であるとＡｌ（ＮＣ）として析出しないため、ＮとＣの供給源としては作用せず、２．０％を超えると鋼材の表面性状を悪化させ疲労特性を悪化させるので、その適正添加範囲を０．００２％〜２．０％とした。
【００３４】
Ｎ：Ｎは鋼の歪み時効硬化の発現およびミクロ組織の制御に必須の添加元素である。しかし、０．０００５％未満であると未変形部分における塗装焼付処理前後の降伏強度の変化量（ＹＳ＿ＢＨ０）＝３０ＭＰａ以上が達成できない。また、０．２％を超えると耐衝撃特性が急激に悪化し、また溶接性も低下する。このため本発明ではＮの範囲を０．０００５〜０．２％に限定した。
【００３５】
Ｃｕ：Ｃｕは塗装焼付工程中にフェライト粒内に析出する炭窒化物、あるいは炭素および窒素からなるクラスター（集合体）の核生成サイトを与えると考えられ、本発明において重要な構成元素の一つである。Ｃ量とＮ量の和に対して０．１倍以上の添加で効果が認められるが、含有量で２．０％を超えるとＣｕの熱間脆性による鋼材の表面割れが顕著になり、また添加量が０．００１％未満であると上　記クラスター核生成サイトを与えない。従って、Ｃｕ含有量の範囲を０．００１〜２．０％の範囲に、またＣｕ量の適正値を（Ｃ量＋Ｎ量）の０．１倍以上に制限した。なお、塗装焼付工程中に安定的により多くの鉄炭窒化物・クラスターの核生成サイトを与えるという観点からは、Ｃｕ含有量は（Ｃ含有量＋Ｎ含有量）の等量（１倍）以上にするのが好適である。
【００３６】
本発明では、上記組成に加えて、Ｎｉの適正量を含有するのが好ましい。
Ｎｉ：ＮｉはＣｕ添加に起因する熱間脆性の抑制とミクロ組織の制御に用いられるので０．１％以上含有することが好ましい。一般的に、添加Ｃｕ量と等量のＮｉを添加するとＣｕによる顕著な熱間割れを抑制できる。従って、その適正添加範囲を２．０％以下に限定した。
【００３７】
本発明では、上記した組成に加えて、更にａ群〜ｄ群のうちの１群または２群以上を含有しても、本発明の目的を達成することができる。
ａ群：Ｃｒ、Ｍｏ、Ｗの１種または２種以上の合計を０．１〜２．０％。
Ｃｒ、Ｍｏ、Ｗは炭窒化物形成元素であり、これらの元素の合計を０．１％以上含有することにより、熱間圧延中、冷却中、あるいは再加熱中に主に炭窒化物として析出させることで、鋼材の強度を調整するのに用いられる。またフェライト粒内に炭窒化物として分布することで、塗装焼付熱処理中に鋼中に再析出するＣおよびＮの供給源として作用する。しかしながら、合計で２．０％を超えると成型加工性の劣化を招く。従って、その合計量の範囲を０．１〜２．０％とした。
【００３８】
ｂ群：Ｎｂ、Ｔｉ、Ｖ、Ｔａのうち１種または２種以上の合計を０．００３〜０．２％。
Ｎｂ、Ｔｉ、Ｖ、Ｔａは炭窒化物形成元素であり、鋼材のミクロ組織およびＣ量、Ｎ量を調整するのに用いられるので、１種又は２種以上の合計を０．００３％以上含有することが好ましい。またフェライト粒内に炭窒化物として分布することで、塗装焼付熱処理中に鋼中に再析出するＣおよびＮの供給源として作用する。しかしながら、合計で０．２％を超えると、成型加工性の劣化を招く。従って、その合計量の範囲を０．２％以下とした。
【００３９】
ｃ群：Ｂを０．０００３〜０．０１０％。
Ｂは０．０００３％以上含有することにより粒界に偏析し、Ｐによる２次加工割れを抑制する効果があり、さらに成形加工性を改善させる効果がある。しかし、０．０１０％を超えると粒界に粗大析出物を形成して、加工割れが発生する。従って、その範囲を０．０００３〜０．０１０％と限定した。
【００４０】
ｄ群：ＣａまたはＭｇのうち１種または２種を合計で０．００１〜０．０１％。
ＣａおよびＭｇは介在物の形態、分布の制御に用いる元素であり、１種又は２種以上を合計で０．００１％以上含有することが好ましい。しかしながら、合計の含有量が０．０１％を超えると、成型加工性の悪化の原因となる。そのため、合計量の範囲を０．００１〜０．０１％とした。
【００４１】
本発明の熱延鋼材では、上記した成分以外の残部はＦｅおよび不可避的不純物である。不可避的不純物として、Ｏは非金属介在物を形成し品質に悪影響を及ぼすので、０．０１％以下にするのが望ましい。
【００４２】
本発明に係る熱延鋼材は塗装焼付工程においてフェライト中に微細炭窒化物あるいは炭素窒素集合体を形成させることが必要であるため、フェライト面積率を５０％以上、好ましくは８０％以上とする。フェライト面積率が１００％でも本発明の効果を奏することができる。残部組織はマルテンサイト、オーステナイト、ラス状ベイナイト、パーライトの１種又は２種以上を含有しても良い。
本発明の組織とするためには、本発明の範囲内にあるように焼き入れ性の小さい成分系とし、熱間加工プロセスにおいて８００℃と５００℃間の平均冷却速度を０．０５〜３０℃／ｓにすることにより得ることができる。
【００４３】
フェライト面積率は圧延方向に平行する断面（Ｌ断面）について、ナイタール液を用いて組織を現出し、次いで光学顕微鏡を用いてミクロ組織を観察した際の明部をフェライト組織と定義し、その部分の面積率を画像解析装置により求める。
【００４４】
次に、フェライト粒内における炭素・窒素集合体、炭窒化物析出物の合計の分布密度については、分布密度がμｍ^３あたり１００個未満であるとＹＳ＿ＢＨ０＝３０ＭＰａ以上の降伏強度上昇が得られなかった。一方、１００００個を超えるような材料は成形加工前の鋼材の常温非時効性が極めて悪いという特徴を有していた。従って、その範囲を体積密度で１００〜１００００μｍ^−３以上に限定した。
【００４５】
なお、クラスターあるいは析出物分布の定量方法としては、直径で約５ｎｍ以上の析出物についてはその大きさに応じて走査型分析電子顕微鏡あるいは透過型分析電子顕微鏡での定量が好適であるが、それ以下の大きさのクラスター、析出物については３次元アトムプローブ電界イオン顕微鏡法（３ＤＡＰ）による定量が好適である。３ＤＡＰによるクラスター有無の判定方法としては、データを３次元原子マップで表示した時に１ｎｍ^３の体積の中に炭素原子と窒素原子が合計で５個以上含む領域をクラスターと判断する方法が簡易で好適である。
【００４６】
なお、炭素・窒素原子集合体とは、炭素のみから構成されていても、窒素のみから構成されていても、炭素と窒素の両方から構成されていても、その強度に及ぼす効果に大きな差違はない。炭窒化物の種類としては、Ｆｅが主成分のものがＣ、Ｎの供給源として好適であるが、Ａｌ、Ｃｒ、Ｍｏ、Ｗ、Ｎｂ、Ｔｉ、Ｖ、Ｔａのうち一種が主成分のものでも同様の供給作用を有する。また炭化物でも窒化物でも炭窒化物の何れでも同様の効果が得られる。
【００４７】
次に、焼付硬化量の差の限定理由について説明する。未変形部分における塗装焼付処理前後の降伏強度変化量ＹＳ＿ＢＨ０が３０ＭＰａ未満であると成形部材の衝撃圧壊変形時に不安定変形を示し衝撃吸収エネルギーも低下する傾向にある。また材料成形加工後において２％変形部分の焼付硬化量と未変形部分の焼付硬化量の差がＹＳ＿ΔＢＨで９０ＭＰａを超えるかあるいはＴＳ＿ΔＢＨで５０ＭＰａを超えると、同様に成形部材の衝撃変形時に不安定変形を示し衝撃吸収エネルギーも低下する。従って、未変形部分における塗装焼付処理前後の降伏強度変化量ＹＳ＿ＢＨ０は３０ＭＰａ以上に制限し、材料成形加工後において変形が加わった部分の焼付硬化量と未変形部分の焼付硬化量の差については、ＹＳで９０ＭＰａ以下かつＴＳで５０ＭＰａ以下に限定した。
【００４８】
一方、未変形部分における塗装焼付処理前後の降伏強度変化量ＹＳ＿ＢＨ０の上限は特に定めることなく本発明の効果を奏することができるが、常温非時効性確保のためには、１５０ＭＰａ以下とすることが好ましい。また、材料成形加工後において変形が加わった部分と未変形部分の降伏強度ＹＳ、引張強度ＴＳの差の下限については、特に定めることなく本発明の効果を奏することができ、その差の値は３０ＭＰａであることがより好ましい。
【００４９】
次に、製造方法の限定理由について説明する。
スラブ加熱温度が１０００℃未満では、鋳造インゴット中に存在する炭窒化物が十分に溶解せず、その後の工程における組織制御が難しくなる。また１０００℃未満では、鋳造マクロ偏析が残留するため、鋼材組織が不均一化し、優れた衝撃特性あるいは疲労特性を達成することが難しい。そのため、その加熱温度の範囲を１０００℃以上に限定した。なお、炭窒化物の溶解をより進行させ、Ａ１　温度以上の最終仕上圧延の温度を確保するために、１１００℃以上の加熱がより好ましい。一方、スラブ加熱中の鋼材酸化防止のためには、加熱温度を１２７０℃以下とする必要がある。
【００５０】
スラブ加熱後の粗圧延の条件は特に規定する必要はなく、常法に従って行えばよい。続いて行う仕上げ圧延については、圧延パスの回数、各パスの温度・圧下率については特に規定する必要はなく常法に従い、成分に応じて所望のミクロ組織を得るように条件を選択すればよい。しかし、本発明においては析出した炭窒化物に圧延を加えて析出物を不安定化させ、電着焼付塗装熱処理中に再析出する炭素・窒素集合体あるいは炭窒化物のＣ、Ｎ供給源とさせることが重要な条件の一つであるため、多くの炭窒化物がまだ析出しない１０００℃を超える温度で仕上げ圧延を終了してしまうと、未変形部分において３０ＭＰａ以上の降伏強度上昇が得られない場合がある。従って、最終仕上げ圧延温度については１０００℃以下の範囲に制限した。一方、鋼材の生産性を低下させないためには、最終仕上げ圧延温度はＡ１　温度以上とする必要がある。
【００５１】
なお、熱間加工時に潤滑圧延を行うことは材質の均質化を図る上で有効である。その場合、摩擦係数は０．１〜０．５の範囲が好ましい。材質の不均一を少なくすることは疲労特性および耐衝撃性向上の観点から好ましい。
【００５２】
仕上げ圧延後の冷却方法については特に規定する必要が無く、水冷、空冷およびこの２種の組み合わせにより行えばよい。ただし仕上げ圧延後に巻取処理を行わない場合には、Ａ１　温度〜１００℃間を平均値で３０℃／ｓを超える速度で冷却を行うと、ＣあるいはＮ供給源になりうる微細炭窒化物の析出が起こらず、その結果、未変形部分において３０ＭＰａ以上のＢＨ量が得られない場合がある。従って、巻取処理を行わない場合のＡ１　温度〜１００℃間の平均冷却速度を３０℃／ｓ以下に制限した。一方、粗大炭窒化物形成の抑制および微細分散化の達成のために、平均冷却速度を０．０１℃／ｓ以上とする。
【００５３】
巻取工程を行う場合は、１００℃未満であると電着焼付塗装熱処理中に再析出する炭素・窒素集合体あるいは炭窒化物のＣ、Ｎ供給源になりうる炭窒化物が十分に析出しないため、巻取温度の範囲を１００℃以上に制限した。一方、粗大炭窒化物形成の抑制および微細分散化達成のために、巻取温度をＡ１　温度以下とした。巻取工程中においてフェライト粒内に微細に炭窒化物を析出させておくほど、塗装焼付工程時に未変形部でより大きな降伏強度上昇（ＹＳ＿ＢＨ０）が得られるので、巻取温度は５５０℃以下で行うことがより好適である。
【００５４】
冷却の後、再加熱処理を行う目的は、フェライト粒内にＣあるいはＮ供給源になりうる炭窒化物を析出させるためである。１００℃未満であると炭窒化物の析出がおこらず、またＡ１　温度を超えるとその前工程で析出していた炭窒化物の溶解が始まってしまう。従って、その温度範囲を１００℃〜Ａ１　温度に制限した。再加熱工程中においてフェライト粒内に微細な炭窒化物を高い分布密度で析出させておくほど、塗装焼付工程時に未変形部でより大きな降伏強度上昇（ＹＳ＿ＢＨ０）が得られるので、再加熱温度は５５０℃以下で行うことがより好適である。
【００５５】
本発明では、析出させた炭窒化物をより不安定化させるために、一連の熱延工程に続いてさらに伸び率０．５〜４０％の調質圧延、またはレベラー加工または矯直加工を施すことがより好ましい。調質圧延またはレベラー加工を施すことにより、未変形部分のΔＢＨ量をより安定的に向上させることができる。なお、析出物が不安定化するとは、析出物が加工時に導入される加工転位と相互作用することによって、析出物の構造中に欠陥が導入されエネルギー的に不安定化すること、あるいは析出物周囲に欠陥が導入され界面エネルギーが上昇しその結果析出物がエネルギー的に分解されやすい状態になることを意味する。
【００５６】
未変形部分のΔＢＨ量向上という観点からは調質圧延量が大きい方がより好ましい。伸び率が０．５％未満であると析出物を不安定化させる効果は小さく、また４０％を超えると鋼材の延性が低下する。従って、調質圧延における伸び率の範囲を０．５〜４０％の範囲に制限した。
【００５７】
【実施例】
次にこの発明を実施例により詳細に説明する。
表１に示す成分のスラブを粗圧延で３０ｍｍ厚にした後に、表２に示す条件で加工熱処理をした。このようにして得られた鋼材について、引張試験、ＢＨ試験、衝撃試験、疲労試験および組織観察を行った。各試験、観察の条件を以下に示す。
【００５８】
引張試験はＪＩＳ１３Ｂ試験片を用い、歪み速度１０^−３／ｓの条件で行った。ＢＨ試験は、次の２つの試験を行った。第１の試験は未変形部分の特性を評価するためであり、１７０℃×２０分間の塗装焼付処理相当の時効処理を施した後、引張試験を行い、降伏強度および引張強度を測定した。ついで、この降伏強度および引張強度と、熱延ままの材料の降伏強度と引張強度の差、すなわちＹＳ＿ＢＨ０とＴＳ＿ＢＨ０（図１参照）を求めた。
【００５９】
第２の試験は変形部分の特性を評価するためであり、２％の引張予歪みを付加した後、一旦除荷し、１７０℃×２０分間の塗装焼付処理相当の時効処理を施した後、引張試験を行い、降伏強度および引張強度を測定した。ついで、この降伏強度および引張強度と、熱延ままの材料の降伏強度と引張強度の差、すなわちＹＳ＿ＢＨ２とＴＳ＿ＢＨ２（図１参照）を求めた。最後に、ＹＳ＿ＢＨ２とＹＳ＿ＢＨ０の差を求め、変形部分と未変形部分の降伏強度上昇量の差ＹＳ＿ΔＢＨとした。また同様にして、ＴＳ＿ＢＨ２とＴＳ＿ＢＨ０の差を求め、それを変形部分と未変形部分の引張強度上昇量の差ＴＳ＿ΔＢＨとした。
【００６０】
耐衝撃特性は、歪み速度１０^３／ｓの高速引張を行い、応力−歪み曲線図を作成した後、歪み量２０％までの積分値を求め、吸収エネルギーを求めることにより評価した。１７０℃×２０分間の塗装焼付処理相当の時効処理を施したもの　　（未変形部分に相当）の吸収エネルギーＥ０と、２％の引張予歪みを付加した後、一旦除荷し１７０℃×２０分間の塗装焼付処理相当の時効処理を施したもの（変形部分に相当）の吸収エネルギーＥ２を求め、さらにその比Ｅ２／Ｅ０を算出し、衝突変形特性と変形の安定性の指標とした。この値が小さいほど、変形部と未変形部の特性差が小さいことを意味し、部材にしたときに安定的に良好な衝撃圧壊特性が得られることを示している。
【００６１】
疲労試験は、熱延鋼板から疲労試験片を採取した後、１７０℃×２０分間の塗装焼付処理相当の時効処理を施したものと、２％の引張予歪みを付加した後一旦除荷し、１７０℃×２０分間の塗装焼付処理相当の時効処理を施したものを用い、疲労試験を行った。Ｓ−Ｎ曲線からそれぞれの疲労限（２×１０^６回）を求め、次いで、両者の疲労限の差と１７０℃×２０分間の塗装焼付処理相当の時効処理を施した材料の疲労限の比を算出し、疲労特性の指標とした。この値が小さいほど、変形部と未変形部の特性差が小さいことを意味し、部材にしたときに安定的に良好な疲労特性が得られることを示している。
【００６２】
１７０℃×２０分間の塗装焼付処理相当の時効処理を施した試験材のフェライト粒内のクラスターあるいは析出物分布は直径で約５ｎｍ以上の析出物については透過型分析電子顕微鏡で観察を行い、それ以下の大きさのクラスター、析出物についてはアトムプローブ電界イオン顕微鏡法によりその分布の観察を行った。３ＤＡＰにおいては、データを３次元原子マップで表示した時に１ｎｍ^３の体積の中に炭素原子と窒素原子が合計で５個以上含む領域をクラスターと判断し、クラスターあるいは析出物の分布を体積密度として求めた。これらの試験結果を表３に示す。
【００６３】
表３に示すように、変形が加わった部分の焼付硬化量と未変形部分の焼付硬化量の差ＹＳ＿ΔＢＨが９０ＭＰａを超えるか、ＴＳ＿ΔＢＨで５０ＭＰａを超えるか、あるいは未変形部分における塗装焼付処理前後の降伏強度変化量（ＹＳ＿ＢＨ０）が３０ＭＰａ未満である場合には、変形部分と未変形部分の耐衝撃特性の差が大きくなり（すなわち吸収エネルギー比が大きくなり）、実部材としての衝撃特性は劣化する傾向になる。実際に、Ｅ２／Ｅ０と実部材の圧壊形態の間には相関関係があり、Ｅ２／Ｅ０が１．２を超えると角柱の高速圧壊試験において倒壊等の不安定圧壊を示し、耐衝撃特性が悪化していた。同様の傾向が疲労特性についてほぼ当て嵌まる。
【００６４】
鋼材Ｎｏ．２は鋼材のＡｌ量が発明範囲以下であるために未変形部分において十分なＹＳ＿ＢＨ０が得られなかった例である。鋼板Ｎｏ．５、Ｎｏ．８、Ｎｏ．２０、Ｎｏ．２２、Ｎｏ．２４はＣｕの添加量が（Ｃ＋Ｎ）量に対して十分でなかったために、十分なＹＳ＿ＢＨ０が得られないか、ＹＳ＿ΔＢＨが大きくなりすぎた例である。また鋼板Ｎｏ．１７は仕上げ圧延後においてＡ１　→１００℃間の平均冷却速度が本発明範囲よりも大きかったためフェライト粒内で炭窒化物の析出が十分に起こらず、その結果十分なＹＳ＿ＢＨ０が得られなかった例である。
【００６５】
【表１】

【００６６】
【表２】

【００６７】
【表３】

【００６８】
【表４】

【００６９】
【発明の効果】
本発明は、電着塗装焼付処理を施す自動車用の構造部材・足廻り部材・パネル部材用途、電機製品用内外板パネル、建築物等の構造物用途に好適な、良好な加工性を有しかつ実部材として使用したときに優れた衝撃特性を呈する歪時効硬化型熱延鋼材素材を安価に提供することができ工業的に価値が高い。さらに優れた耐衝撃性に加えて、優れた疲労特性を併せ持つ部材を提供できることから、自動車や構造物の安全性向上に格段の効果を有する。
【図面の簡単な説明】
【図１】ＹＳ＿ＢＨ０値、ＹＳ＿ΔＢＨ値、ＴＳ＿ΔＢＨ値の定義の模式図[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a strain age hardening type hot rolled steel structural member excellent in impact characteristics and molding workability suitable for use in structural members, undercarriage members, and panel members of an automobile, and a method of manufacturing the hot rolled steel material. It is applicable to a wide range of steel materials having a tensile strength of about 250 MPa to 900 MPa. Since the present invention also has excellent fatigue properties, it can be applied to tire wheel applications and structural applications such as buildings, ships, and bridges in addition to the above applications.
[0002]
[Prior art]
CO emissions from automobiles from the viewpoint of global environmental protection ₂ Mitigation is an important issue. CO ₂ To reduce the weight, it is effective to reduce the weight of the vehicle body. Therefore, there is an increasing need for higher strength steel materials. However, it is generally known that increasing the strength of a material is accompanied by deterioration in press formability, such as a decrease in shape freezing property and cracking during molding. Was desired. Other characteristics required for automotive steel are properties that do not cause material deterioration due to vibration during traveling, that is, fatigue properties, and materials for absorbing energy on the material side during high-speed collisions and ensuring occupant safety. Characteristics, that is, characteristics such as impact resistance.
[0003]
As for structural steel materials used for buildings, bridges and ships, there is a high demand for high-strength steel materials for the purpose of reducing the weight of structures. As in the case of the above-mentioned steel materials for automobiles, there is a strong demand for molding workability and impact resistance and fatigue properties on the premise of collision between structures.
[0004]
In response to such demands, as a technique for achieving high strength while ensuring moldability, it is kept soft during molding and occurs in an electrodeposition coating baking process at 170 ° C. × 20 minutes after molding. There is known a technique using so-called bake hardenability (BH) in which the yield strength or tensile strength is increased by using the strain age hardening phenomenon. In this type of steel, carbon or nitrogen atoms are dissolved during forming to secure formability, and carbon or nitrogen atoms are fixed to dislocations generated in the steel during forming in the electrodeposition coating baking process. In this case, the yield strength or tensile strength is increased by dispersing or precipitating carbides or nitrides on dislocations.
[0005]
On the other hand, regarding the impact characteristics, it is thought that the higher the deformation stress at the deformation at a high strain rate, the more energy is absorbed during the impact deformation, and the impact resistance is improved by utilizing the strain age hardening phenomenon. A technique for causing this to occur has been proposed. For example, Patent Literature 1 below discloses that in a given steel component, by controlling the amounts of C and N in a solid solution state and the ferrite crystal grain size, both the yield stress and the tensile strength are increased and the impact resistance is excellent. A strain age hardening type hot rolled steel sheet and a manufacturing method thereof have been proposed. Patent Literature 2 below proposes a strain-age hardening type high-tensile hot-rolled steel sheet which is excellent in impact resistance and uses a solid solution of N, in which the ferrite grain size is controlled, and a method of manufacturing the same.
[0006]
[Patent Document 1]
JP 2001-335889 A
[0007]
[Patent Document 2]
JP 2001-226744 A
[0008]
However, members are often formed by bending, pressing, or the like. However, in actual members, unformed portions (undeformed portions) occupy a considerable proportion of the member volume. No matter how much the deformed part is strengthened by the strain hardening phenomenon, the strength of the soft undeformed part will ultimately determine the impact resistance. (2) The strain age hardening type proposed so far There was a problem that the method of impact destruction when impact deformation was applied was not reproducible even if steel material was formed → electrodeposition coating and baked, so it was practically difficult to apply it to actual members. .
[0009]
[Problems to be solved by the invention]
The present invention has been made in view of the above-mentioned situation, and is used for structural members, undercarriage members, and panel members for automobiles made through an electrodeposition coating process, structural members for buildings, and inner and outer plates of electric products. An object of the present invention is to provide a strain-age hardening type hot-rolled steel structural member having good workability and excellent impact properties, which is suitable for a panel, and a method for producing the hot-rolled steel material.
[0010]
[Means for Solving the Problems]
The present inventors have conducted a high-speed compression deformation test of a real member to solve the above-mentioned problem. As a result, the cause of unstable deformation during impact crushing deformation is as follows. The reason why it is difficult to obtain further excellent impact resistance is that if the difference between the deformation stresses of the deformed part and the undeformed part becomes large, the absorbed energy at the time of high-speed deformation will eventually become the characteristic of the undeformed part. It was decided that it was because it was decided. Further, as a cause of the large non-uniformity of the strength of the structural member after performing the electrodeposition coating baking process, dislocations are introduced in portions where the deformation is locally applied by the molding process, and the strain age hardens. Since dislocations are not introduced in the deformed portion, it was concluded that age hardening did not occur even in the baking coating process, or that the hardening amount was small.
[0011]
Therefore, the present inventors have conducted intensive experiments and studies in order to achieve a strain age hardening type hot rolled steel material having excellent impact resistance when molded into a real structural member, and as a result, after the electrodeposition baking coating process It is necessary to obtain a strength increase within a predetermined range even in the undeformed part, and to reduce the difference in strength between the deformed part and the undeformed part after the electrodeposition baking coating process, that is, to minimize uneven strength as much as possible. It was found that it was effective for improvement.
The present inventors have found that in order to reduce the difference in strength between the deformed portion and the undeformed portion, it is effective to cause clustering and precipitation of carbon or nitrogen atoms even in ferrite grains in the undeformed portion. In addition, it has been clarified that it is effective to increase the Cu / (C + N) component ratio in steel to an appropriate value or more in order to obtain such a structure.
[0012]
As a result of a detailed analysis of the mechanism in which such a phenomenon occurs, it was found that by adding an appropriate amount of Cu, Cu atoms or vacancy aggregates introduced by the addition of Cu become nucleation sites and where dislocations did not exist. However, we have found a completely new fact that fine carbonitrides or carbon-nitrogen clusters precipitate during electrodeposition coating baking.
[0013]
Furthermore, the present inventors have investigated the fine carbonitride and carbon and nitrogen cluster C and N supply sources newly re-precipitated in the grains during the coating baking treatment, and found that (1) solid solution in the grains In addition to the previously proposed sources such as C or N, (2) C or N segregated at the grain boundaries, (3) destabilization present in the ferrite grains before bake coating He revealed for the first time a completely new fact that carbonitride could be a source. Specifically, if the precipitate is destabilized in terms of energy by a hot rolling process, a cold working process, a temper rolling process, or a leveler process, etc. before the forming process, it becomes age hardened in the paint baking process. It has been found for the first time to be a source of contributing C and N.
[0014]
The present inventors have developed a technique for adding an appropriate amount of Cu, which has the effect of finely depositing carbon / nitrogen aggregates or carbonitrides even when processing dislocations do not exist, as described above. By mutually utilizing technologies that use unstable precipitates as a source of carbon and nitrogen to nitride, even in the undeformed area, heat treatment equivalent to electrodeposition coating sufficiently hardens, and the deformed part due to molding processing The present inventors have found that it is possible to reduce the difference in strength between the undeformed portion and the undeformed portion, and have accomplished the present invention.
[0015]
In addition, as a result of the study by the present inventors, the above composition makes it possible to obtain one or more of martensite, retained austenite, and pearlite in addition to the structure of ferrite single phase, in addition to the structure of mainly ferrite. Even in a dual phase steel containing the structure of, a strain-aged hardened hot-rolled steel with excellent impact properties and moldability can be obtained, and in terms of the strength level, the tensile strength is within a wide range of 250 to 900 MPa. It was found that superior impact properties were obtained compared to steel.
[0016]
The present invention takes the following measures in order to solve the above problems. That is,
(1) A first invention is a structural member before hot-rolled steel material subjected to an electrodeposition coating baking treatment, and is a mass%
C: 0.0005 to 0.2%, Si: 0.001 to 2.0%,
Mn: 0.001 to 2.0%, P: 0.2% or less,
S: 0.015% or less, Al: 0.002 to 2.0%,
N: 0.0005 to 0.2%, Cu: 0.001 to 2.0%
And the Cu content is 0.1 times or more of the sum of the C content and the N content, and the balance is composed of Fe and unavoidable impurities. The yield strength change (YS_BH0) before and after the heat treatment in a portion that has not been deformed by the forming process before and after the heat treatment is 30 MPa or more.
[0017]
(2) The second invention is the strain-age hardening type hot-rolled steel structural member according to the above (1), wherein the ferrite area ratio is 50% or more,
(3) The third invention is the invention according to the above (1) or (2), in which the yield strength YS and the tensile strength of a portion deformed by molding and an undeformed portion when the heat treatment is performed. The difference in strength TS is 90 MPa or less for YS and 50 MPa or less for TS, respectively.
[0018]
(4) A fourth invention is the carbon / nitrogen atom aggregate in the undeformed portion ferrite grains when the heat treatment is performed, in addition to the invention described in any one of the above (1) to (3). And the distribution density of the total of the two types of carbonitride precipitates is 100 to 10,000 ^-3 Characterized in that
(5) A fifth invention is characterized in that in addition to the composition according to any one of the above (1) to (4), Ni: 0.1 to 2.0% by mass%.
[0019]
(6) The sixth invention is characterized in that, in addition to the above composition, one or more of the following groups a to d are further included in mass%.
Group a: 0.1 to 2.0% of one or more of Cr, Mo and W
Group b: 0.03 to 0.2% of the total of one or more of Nb, Ti, V, and Ta
Group c: B 0.0003-0.010%
Group d: 0.001 to 0.01% of one or two of Ca or Mg in total
(7) A seventh invention provides strain-age hardening excellent in impact characteristics, characterized in that the hot-rolled steel material according to any one of the above (1) to (6) is subjected to electroplating or hot-dip plating. Die hot rolled steel.
[0020]
Further, the present invention is a method for producing a strain-aged hardened hot-rolled steel excellent in impact properties,
(8) In an eighth aspect, the steel slab having the composition described in the above (1), (5) or (6) is heated to a slab heating temperature of 1000 to 1270 ° C., rough-rolled, and then A1 to 1000. C., finishing the finish rolling at a temperature of A1 and cooling at a mean cooling rate of 0.01 to 30 ° C./s between A1 and 100 ° C. by air cooling, water cooling or a combination of these two cooling methods.
(9) A ninth aspect of the present invention is to heat a steel slab having the composition described in (1), (5) or (6) to a slab heating temperature of 1000 to 1270 ° C., perform rough rolling, and then A1 to 1000 ° C. After finishing the finish rolling, cooling by air cooling, water cooling or a combination of these two cooling methods, and then winding at a temperature between A1 and 100 ° C.
[0021]
(10) In a tenth aspect, a steel slab having the composition described in the above (1), (5) or (6) is heated to a slab heating temperature of 1000 to 1270 ° C., rough-rolled, and then A1 to 1000. After finishing the finish rolling at ℃, and cooling to room temperature by air cooling, water cooling or a combination of these two cooling methods, it is characterized by performing a reheating treatment between 100 ℃ ~ A1 temperature,
(11) The eleventh invention is characterized in that, following the production process described in any one of the above (8) to (10), temper rolling or leveler processing with an elongation percentage of 0.5 to 40% is performed. And
[0022]
In the actual forming process, the amount of prestrain and its deformation mode (tensile or compressive, uniaxial or biaxial tensile, etc.) differ depending on the location of the member, but the present inventors have found that various types of forming As a result of a detailed examination of the relationship between the strain aging hardening behavior and the impact characteristics, it was found that the characteristics of the deformed portion can be well represented by a tensile strain of 2%. Based on this knowledge, in the present invention, the tensile deformation is set to 2% as the pre-deformation amount of the deformed portion and evaluated. The undeformed portion means that no additional strain is applied when the material is formed and formed into a structural member, and deformation due to roll rolling or die rolling added during the production of the steel material is added. No problem. Yield strength refers to the yield strength that is correlated with the impact resistance. FIG. 1 shows a schematic diagram of the definitions of the YS_BH0 value, the YS_ΔBH value, and the TS_ΔBH value.
[0023]
The baking coating conditions were evaluated by a heat treatment of 170 ° C. × 20 minutes which is used as a standard in the examples. However, the baking coating conditions after the forming process in which the present invention is effective should not be limited to 170 ° C. × 20 minutes. If YS_BH0 of 30 MPa or more can be obtained, the baking temperature is 170 ° C. or less, for example, 100 ° C. The object of the present invention can be achieved even at a low-temperature baking condition of ℃ and a baking time of 20 minutes or less.
[0024]
In the present invention, the “carbon / nitrogen atom aggregate” is an atomic aggregate composed of carbon or nitrogen, and refers to an aggregate of five or more atoms including carbon atoms and nitrogen atoms. The term “ferrite” refers to a polygonal ferrite structure, a pseudopolygonal ferrite structure, or a granular bainitic ferrite structure containing an M / A composite as shown in Non-Patent Document 1 below.
[0025]
[Non-patent document 1]
ISIJ international, 35 (2002), pp. 941-944.
[0026]
Regarding the method for analyzing Cu in steel, it is simple and preferable to use spark discharge emission spectroscopy when the concentration exceeds 0.002%, but when the concentration is less than 0.002%, ICP emission analysis is used. Method or secondary ion mass spectrometry. When ICP emission analysis is used, sol. Cu and insol. It is necessary to measure the total amount of Cu including Cu.
[0027]
The above-described hot-rolled steel material of the present invention is suitable as a raw material for various platings. The plating layer may be formed by either an electroplating method or a hot-dip plating method, and zinc, chromium, tin, and nickel are exemplified as main components of the plating.
Further, the steel of the present invention has not only excellent impact resistance but also excellent fatigue resistance when molded as a structural member.
[0028]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described in detail.
First, the reasons for limiting the components will be described. The component content is% by mass.
C: C is an additive element essential for the development of strain age hardening and control of the microstructure of steel. However, if it is less than 0.0005%, a change in yield strength (YS_BH0) of the undeformed portion before and after the coating baking process cannot be achieved of 30 MPa or more. On the other hand, if it exceeds 0.2%, the impact resistance is rapidly deteriorated, and the weldability is also reduced. Therefore, in the present invention, the range of C is limited to 0.0005 to 0.2%.
[0029]
Si: Since Si is used for adjusting the microstructure and strength of the steel material, it should be contained at 0.001% or more. However, if the content exceeds 2.0% by weight, the descaling property at the time of hot rolling is deteriorated and the fatigue properties are deteriorated. Therefore, the Si content was limited to a range of 2.0% or less.
[0030]
Mn: Since Mn is used for adjusting the microstructure and strength of the steel material, Mn is contained at 0.001% or more. However, if it exceeds 2.0%, the moldability deteriorates. Therefore, the Mn content was limited to a range of 2.0% or less. From the viewpoint of suppressing hot brittleness of S, 0.005% or more is desirable.
[0031]
P: P, which is an impurity, is used for adjusting the strength of the steel material. However, if it exceeds 0.2%, there is a possibility that a crack may occur at the time of forming, so the range of the P content is set to 0.2% or less.
[0032]
S: S, which is an impurity, is present in steel as MnS and CuS, and is used for adjusting the strength and ductility of the steel material by controlling the crystal grain size. The fine distribution within the grains also serves as a supply source of C and N reprecipitated in the steel during the coating baking heat treatment. However, if it exceeds 0.015%, there is a possibility that hot embrittlement may occur. Therefore, the range is limited to 0.015% or less.
[0033]
Al: An important effect of Al is that if it is precipitated as Al (NC) in the ferrite grains before the coating baking treatment, the carbon / nitrogen aggregates or N, C that form carbon nitrides re-precipitate during the coating baking treatment. It acts as a supply source, and as a result, the BH in the undeformed region increases by simultaneous addition with Cu. However, if the added amount of Al is less than 0.002%, it does not precipitate as Al (NC), so it does not act as a supply source of N and C. If it exceeds 2.0%, the surface properties of the steel material deteriorate. Since the fatigue characteristics are deteriorated, the proper addition range is set to 0.002% to 2.0%.
[0034]
N: N is an additive element essential for the development of strain age hardening and control of microstructure of steel. However, if it is less than 0.0005%, a change in yield strength (YS_BH0) of the undeformed portion before and after the coating baking process cannot be achieved of 30 MPa or more. On the other hand, if it exceeds 0.2%, the impact resistance is rapidly deteriorated, and the weldability is also reduced. Therefore, in the present invention, the range of N is limited to 0.0005 to 0.2%.
[0035]
Cu: Cu is considered to provide a nucleation site for carbonitride or a cluster (aggregate) composed of carbon and nitrogen which precipitates in ferrite grains during the coating baking step, and is one of the important constituent elements in the present invention. It is. The effect is recognized when the content is 0.1 times or more of the sum of the C content and the N content. However, when the content exceeds 2.0%, the surface cracks of the steel material due to hot embrittlement of Cu become remarkable, and If the addition amount is less than 0.001%, the above-mentioned cluster nucleation site is not provided. Therefore, the range of the Cu content is limited to the range of 0.001 to 2.0%, and the appropriate value of the Cu amount is limited to 0.1 times or more of (C amount + N amount). In addition, from the viewpoint of stably providing more nucleation sites for iron carbonitrides / clusters during the coating baking process, the Cu content is equal to or more than (equal to 1 times) the (C content + N content). It is preferred to do so.
[0036]
In the present invention, it is preferable to contain an appropriate amount of Ni in addition to the above composition.
Ni: Ni is preferably used in an amount of 0.1% or more because it is used for suppressing hot brittleness and controlling the microstructure due to the addition of Cu. Generally, when Ni is added in an amount equal to the amount of added Cu, remarkable hot cracking due to Cu can be suppressed. Therefore, the appropriate addition range is limited to 2.0% or less.
[0037]
In the present invention, the object of the present invention can be achieved by further containing one or more of the groups a to d in addition to the above composition.
Group a: 0.1 to 2.0% of one or more of Cr, Mo and W.
Cr, Mo, and W are carbonitride forming elements. By containing 0.1% or more of the total of these elements, they are mainly precipitated as carbonitrides during hot rolling, cooling, or reheating. This is used for adjusting the strength of the steel material. In addition, by distributing as carbonitride in ferrite grains, it acts as a supply source of C and N re-precipitated in steel during paint baking heat treatment. However, when the total exceeds 2.0%, the moldability deteriorates. Therefore, the range of the total amount is set to 0.1 to 2.0%.
[0038]
Group b: 0.003 to 0.2% of one or more of Nb, Ti, V, and Ta.
Nb, Ti, V, and Ta are carbonitride forming elements and are used to adjust the microstructure and the C and N contents of the steel material. Therefore, the total content of one or two or more elements is 0.003% or more. Is preferred. In addition, by distributing as carbonitride in ferrite grains, it acts as a supply source of C and N re-precipitated in steel during paint baking heat treatment. However, if the total exceeds 0.2%, the moldability deteriorates. Therefore, the range of the total amount is set to 0.2% or less.
[0039]
Group c: B 0.0003 to 0.010%.
When B is contained in an amount of 0.0003% or more, B segregates at the grain boundary, has an effect of suppressing secondary work cracking due to P, and has an effect of further improving the formability. However, if it exceeds 0.010%, coarse precipitates are formed at the grain boundaries, and work cracks occur. Therefore, the range was limited to 0.0003 to 0.010%.
[0040]
Group d: 0.001 to 0.01% of one or two of Ca and Mg in total.
Ca and Mg are elements used for controlling the form and distribution of inclusions, and it is preferable that one or more of them contain 0.001% or more in total. However, when the total content exceeds 0.01%, it causes deterioration in moldability. Therefore, the range of the total amount is set to 0.001 to 0.01%.
[0041]
In the hot-rolled steel material of the present invention, the balance other than the above components is Fe and inevitable impurities. As an unavoidable impurity, O forms nonmetallic inclusions and adversely affects the quality.
[0042]
In the hot-rolled steel material according to the present invention, it is necessary to form fine carbonitrides or carbon-nitrogen aggregates in the ferrite in the coating baking step, so that the ferrite area ratio is 50% or more, preferably 80% or more. Even when the ferrite area ratio is 100%, the effects of the present invention can be obtained. The remaining structure may contain one or more of martensite, austenite, lath bainite, and pearlite.
In order to obtain the structure of the present invention, a component system having a small hardenability is used as in the range of the present invention, and the average cooling rate between 800 ° C and 500 ° C in the hot working process is 0.05 to 30 ° C. / S.
[0043]
For the ferrite area ratio, a cross section parallel to the rolling direction (L cross section) reveals a microstructure using a nital solution, and then a bright portion when the microstructure is observed using an optical microscope is defined as a ferrite microstructure. Is determined by an image analyzer.
[0044]
Next, regarding the total distribution density of the carbon / nitrogen aggregates and carbonitride precipitates in the ferrite grains, the distribution density was μm ³ If less than 100 pieces per piece, a yield strength increase of YS_BH0 = 30 MPa or more could not be obtained. On the other hand, a material having more than 10,000 pieces had a characteristic that the non-aging property at room temperature of the steel material before forming was extremely poor. Therefore, the range is 100 to 10000 μm in volume density. ^-3 Limited to the above.
[0045]
As a method for quantifying the distribution of clusters or precipitates, for a precipitate having a diameter of about 5 nm or more, quantification using a scanning analytical electron microscope or a transmission analytical electron microscope is preferable according to the size. For clusters and precipitates having the following sizes, quantification by three-dimensional atom probe field ion microscopy (3DAP) is preferable. As a method of determining the presence or absence of a cluster by 3DAP, when the data is displayed in a three-dimensional atomic map, 1 nm ³ It is simple and preferable to determine a region that includes a total of 5 or more carbon atoms and nitrogen atoms in the volume of a cluster as a cluster.
[0046]
In addition, the carbon / nitrogen atom aggregate is a significant difference in its effect on strength whether it is composed only of carbon, composed only of nitrogen, or composed of both carbon and nitrogen. Absent. As for the type of carbonitride, those mainly composed of Fe are suitable as the supply source of C and N, but those of one of Al, Cr, Mo, W, Nb, Ti, V and Ta are mainly used. However, it has a similar supply action. Similar effects can be obtained with any of carbide, nitride and carbonitride.
[0047]
Next, the reason for limiting the difference in the bake hardening amount will be described. If the yield strength change YS_BH0 of the undeformed portion before and after the coating baking treatment is less than 30 MPa, the molded member will exhibit unstable deformation during impact crushing deformation, and the impact absorption energy tends to decrease. Similarly, if the difference between the baking hardening amount of the 2% deformed portion and the baking hardening amount of the undeformed portion exceeds 90 MPa in YS_ΔBH or exceeds 50 MPa in TS_ΔBH after the material forming process, similarly, unstable deformation occurs upon impact deformation of the formed member. And the shock absorption energy also decreases. Therefore, the yield strength change YS_BH0 before and after the coating baking process in the undeformed portion is limited to 30 MPa or more, and the difference between the baking hardening amount of the deformed portion and the baking hardening amount of the undeformed portion after the material forming process is as follows: It was limited to 90 MPa or less in YS and 50 MPa or less in TS.
[0048]
On the other hand, the effect of the present invention can be exerted without particularly defining the upper limit of the yield strength change amount YS_BH0 before and after the coating baking process in the undeformed portion. However, in order to ensure the normal temperature non-aging property, the upper limit is 150 MPa or less. preferable. In addition, the lower limit of the difference between the yield strength YS and the tensile strength TS of the deformed portion and the undeformed portion after the material forming process can exert the effect of the present invention without any particular determination, and the value of the difference is More preferably, it is 30 MPa.
[0049]
Next, the reasons for limiting the manufacturing method will be described.
If the slab heating temperature is less than 1000 ° C., carbonitrides present in the cast ingot are not sufficiently dissolved, and it becomes difficult to control the structure in the subsequent steps. If the temperature is lower than 1000 ° C., since cast macro segregation remains, the structure of the steel material becomes non-uniform, and it is difficult to achieve excellent impact characteristics or fatigue characteristics. Therefore, the range of the heating temperature is limited to 1000 ° C. or higher. Heating at 1100 ° C. or more is more preferable in order to further advance the dissolution of the carbonitride and secure the final finish rolling temperature at A1 temperature or more. On the other hand, in order to prevent steel material oxidation during slab heating, the heating temperature needs to be 1270 ° C. or less.
[0050]
The conditions of the rough rolling after the slab heating need not be particularly specified, and may be performed according to a conventional method. For the subsequent finish rolling, the number of rolling passes, the temperature and rolling reduction of each pass do not need to be particularly specified, and may be selected according to a conventional method, so as to obtain a desired microstructure according to the components. . However, in the present invention, the precipitated carbonitride is rolled to destabilize the precipitate, and a carbon / nitrogen aggregate or carbonitride C, N supply source reprecipitated during the electrodeposition baking coating heat treatment is used. Since it is one of the important conditions, if the finish rolling is completed at a temperature exceeding 1000 ° C. at which much carbonitride does not yet precipitate, a yield strength increase of 30 MPa or more is obtained in the undeformed portion. May not be. Therefore, the final finish rolling temperature was limited to a range of 1000 ° C. or less. On the other hand, in order not to lower the productivity of the steel material, the final finish rolling temperature needs to be equal to or higher than the A1 temperature.
[0051]
Performing lubrication rolling during hot working is effective in achieving homogenization of the material. In that case, the coefficient of friction is preferably in the range of 0.1 to 0.5. It is preferable to reduce the non-uniformity of the material from the viewpoint of improving fatigue characteristics and impact resistance.
[0052]
There is no need to particularly define the cooling method after the finish rolling, and the cooling method may be water cooling, air cooling, or a combination of the two. However, if the winding process is not performed after the finish rolling, cooling at a rate exceeding 30 ° C./s on average between the temperature of A1 and 100 ° C. can reduce the fine carbonitride that can be a C or N supply source. Precipitation does not occur, and as a result, a BH amount of 30 MPa or more may not be obtained in the undeformed portion. Therefore, the average cooling rate between the A1 temperature and 100 ° C when the winding process is not performed is limited to 30 ° C / s or less. On the other hand, in order to suppress formation of coarse carbonitrides and achieve fine dispersion, the average cooling rate is set to 0.01 ° C./s or more.
[0053]
In the case of performing the winding step, if the temperature is lower than 100 ° C., carbon / nitride aggregates re-precipitated during the heat treatment for electrodeposition baking coating or carbonitrides which can be a C, N supply source of carbonitrides are not sufficiently deposited. Therefore, the range of the winding temperature is limited to 100 ° C. or higher. On the other hand, in order to suppress the formation of coarse carbonitrides and achieve fine dispersion, the winding temperature was set to the A1 temperature or lower. The higher the precipitation of carbonitrides in the ferrite grains during the winding process, the larger the yield strength (YS_BH0) is obtained in the undeformed portion during the coating baking process. It is more preferable to do so.
[0054]
The purpose of performing the reheating treatment after cooling is to precipitate carbonitrides that can serve as a C or N supply source in the ferrite grains. If the temperature is lower than 100 ° C., precipitation of carbonitride does not occur, and if the temperature exceeds A1, the dissolution of carbonitride precipitated in the preceding step starts. Therefore, its temperature range was limited to 100 ° C. to A1 temperature. The higher the precipitation density of fine carbonitrides in the ferrite grains during the reheating step, the higher the yield strength (YS_BH0) is obtained in the undeformed portion during the coating baking step. It is more preferable to carry out at 550 ° C. or lower.
[0055]
In the present invention, in order to further destabilize the precipitated carbonitride, temper rolling with an elongation of 0.5 to 40% or leveler processing or straightening is performed after a series of hot rolling steps. Is more preferable. By performing temper rolling or leveler processing, the ΔBH amount of the undeformed portion can be more stably improved. In addition, the instability of the precipitate means that the precipitate interacts with the processing dislocation introduced during processing, thereby introducing defects in the structure of the precipitate and destabilizing the energy, or This means that defects are introduced into the surroundings, the interface energy increases, and as a result, the precipitates are easily decomposed energetically.
[0056]
From the viewpoint of improving the ΔBH amount of the undeformed portion, it is more preferable that the temper rolling amount is large. If the elongation is less than 0.5%, the effect of destabilizing the precipitate is small, and if it exceeds 40%, the ductility of the steel material is reduced. Therefore, the range of the elongation in the temper rolling was limited to the range of 0.5 to 40%.
[0057]
【Example】
Next, the present invention will be described in detail with reference to examples.
After the slab having the components shown in Table 1 was rough-rolled to a thickness of 30 mm, a working heat treatment was performed under the conditions shown in Table 2. The steel material thus obtained was subjected to a tensile test, a BH test, an impact test, a fatigue test, and a structure observation. The conditions for each test and observation are shown below.
[0058]
The tensile test uses a JIS13B test piece and has a strain rate of 10 ^-3 / S. In the BH test, the following two tests were performed. The first test was to evaluate the characteristics of the undeformed portion. After performing an aging treatment equivalent to a paint baking treatment at 170 ° C. for 20 minutes, a tensile test was performed to measure the yield strength and the tensile strength. Next, the yield strength and the tensile strength, and the difference between the yield strength and the tensile strength of the as-rolled material, that is, YS_BH0 and TS_BH0 (see FIG. 1) were determined.
[0059]
The second test is for evaluating the properties of the deformed portion. After applying a 2% tensile prestrain, the load is temporarily unloaded, and after aging treatment equivalent to a paint baking treatment at 170 ° C. for 20 minutes, A tensile test was performed to measure the yield strength and the tensile strength. Then, the difference between the yield strength and the tensile strength and the difference between the yield strength and the tensile strength of the as-rolled material, that is, YS_BH2 and TS_BH2 (see FIG. 1) were determined. Finally, the difference between YS_BH2 and YS_BH0 was determined, and the difference was YS_ΔBH between the yield strength increase of the deformed portion and the undeformed portion. Similarly, the difference between TS_BH2 and TS_BH0 was obtained, and the difference was used as the difference TS_ΔBH in the amount of increase in tensile strength between the deformed portion and the undeformed portion.
[0060]
The shock resistance is a strain rate of 10 ³ After performing a high-speed tension of / s to create a stress-strain curve diagram, evaluation was performed by obtaining an integrated value up to a strain amount of 20% and obtaining absorbed energy. After applying an aging treatment equivalent to a paint baking treatment at 170 ° C. for 20 minutes (corresponding to an undeformed portion) and adding a 2% tensile pre-strain, the load is once unloaded and 170 ° C. for 20 minutes Absorbed energy E2 of the aging treatment (corresponding to the deformed portion) equivalent to the paint baking treatment was calculated, and the ratio E2 / E0 was calculated, and the ratio was used as an index of collision deformation characteristics and deformation stability. The smaller this value is, the smaller the characteristic difference between the deformed portion and the undeformed portion is, which indicates that a good impact crushing characteristic can be stably obtained when the member is formed.
[0061]
The fatigue test was performed by taking a fatigue test specimen from a hot-rolled steel sheet, subjecting it to an aging treatment equivalent to a paint baking treatment at 170 ° C. for 20 minutes, and once unloading after adding a 2% tensile prestrain, A fatigue test was carried out using an aging treatment equivalent to a paint baking treatment at 170 ° C. for 20 minutes. From the SN curve, each fatigue limit (2 × 10 ⁶ ) Was calculated, and then the ratio of the fatigue limit of the material subjected to the aging treatment equivalent to the paint baking treatment at 170 ° C. for 20 minutes was calculated as the index of the fatigue characteristics. The smaller the value is, the smaller the characteristic difference between the deformed portion and the undeformed portion is, which indicates that a good fatigue characteristic can be obtained stably when the member is formed.
[0062]
The distribution of clusters or precipitates in the ferrite grains of the test material which has been subjected to aging treatment equivalent to paint baking treatment at 170 ° C. for 20 minutes is observed with a transmission analysis electron microscope for precipitates having a diameter of about 5 nm or more. The distribution of clusters and precipitates having the following sizes was observed by atom probe field ion microscopy. In 3DAP, when data is displayed in a three-dimensional atomic map, 1 nm ³ A region containing 5 or more carbon atoms and nitrogen atoms in total in the volume was determined to be a cluster, and the distribution of clusters or precipitates was determined as the volume density. Table 3 shows the test results.
[0063]
As shown in Table 3, the difference YS_ΔBH between the baking hardening amount of the deformed portion and the baking hardening amount of the undeformed portion exceeds 90 MPa, exceeds 50 MPa in TS_ΔBH, or before and after the coating baking process in the undeformed portion. When the yield strength change amount (YS_BH0) is less than 30 MPa, the difference between the impact resistance characteristics of the deformed portion and the undeformed portion increases (that is, the absorbed energy ratio increases), and the impact characteristics as a real member deteriorate. Become a trend. Actually, there is a correlation between E2 / E0 and the crushing form of the actual member, and when E2 / E0 exceeds 1.2, unstable crushing such as collapse occurs in a high-speed crushing test of a prism, and the impact resistance is poor. Was getting worse. A similar trend is almost true for fatigue properties.
[0064]
Steel No. No. 2 is an example in which sufficient YS_BH0 was not obtained in the undeformed portion because the amount of Al in the steel material was below the range of the invention. Steel sheet No. 5, no. 8, No. 20, no. 22, no. 24 is an example in which sufficient YS_BH0 was not obtained or YS_ΔBH became too large because the added amount of Cu was not sufficient with respect to the (C + N) amount. In addition, the steel sheet No. No. 17 is an example in which the average cooling rate between A1 and 100 ° C. after the finish rolling was larger than the range of the present invention, so that precipitation of carbonitride did not sufficiently occur in the ferrite grains, and as a result, sufficient YS_BHO was not obtained. is there.
[0065]
[Table 1]

[0066]
[Table 2]

[0067]
[Table 3]

[0068]
[Table 4]

[0069]
【The invention's effect】
The present invention has a good workability suitable for structural members such as automobile structural members, undercarriage members and panel members to be subjected to electrodeposition coating baking treatment, inner and outer panel panels for electric products, and structural applications such as buildings. In addition, a strain age hardening type hot rolled steel material exhibiting excellent impact characteristics when used as an actual member can be provided at low cost, and is industrially highly valuable. Since it is possible to provide a member having excellent fatigue characteristics in addition to excellent impact resistance, it has a remarkable effect on improving the safety of automobiles and structures.
[Brief description of the drawings]
FIG. 1 is a schematic diagram of definitions of YS_BH0 value, YS_ΔBH value, and TS_ΔBH value.

Claims (11)

熱延鋼材を成形加工した電着塗装焼付処理を行う前の構造部材であって、質量％で、
Ｃ　：０．０００５〜０．２％、
Ｓｉ：０．００１〜２．０％、
Ｍｎ：０．００１〜２．０％、
Ｐ　：０．２％以下、
Ｓ　：０．０１５％以下、
Ａｌ：０．００２〜２．０％、
Ｎ　：０．０００５〜０．２％、
Ｃｕ：０．００１〜２．０％
を含み、かつＣｕ含有量がＣ含有量とＮ含有量の和の０．１倍以上であり、残部がＦｅおよび不可避的不純物からなり、１７０℃×２０分間の電着塗装焼付相当の熱処理を行った際に、成形加工により変形していない部分における熱処理前後の降伏強度変化量（ＹＳ＿ＢＨ０）が３０ＭＰａ以上であることを特徴とする衝撃特性に優れた歪時効硬化型熱延鋼構造部材。It is a structural member before performing the electrodeposition coating baking process formed by processing hot rolled steel material.
C: 0.0005 to 0.2%,
Si: 0.001 to 2.0%,
Mn: 0.001 to 2.0%,
P: 0.2% or less,
S: 0.015% or less,
Al: 0.002 to 2.0%,
N: 0.0005 to 0.2%,
Cu: 0.001 to 2.0%
And the Cu content is 0.1 times or more of the sum of the C content and the N content, and the balance is composed of Fe and unavoidable impurities. A strain age hardening type hot rolled steel structural member excellent in impact characteristics, characterized in that a yield strength change (YS_BH0) before and after heat treatment in a portion which is not deformed by molding when subjected to a forming process is 30 MPa or more. フェライト面積率が５０％以上であることを特徴とする請求項１記載の歪時効硬化型熱延鋼構造部材。2. The strain age hardening type hot rolled steel structural member according to claim 1, wherein the ferrite area ratio is 50% or more. 前記熱処理を行った際に、成形加工により変形した部分と、変形していない部分の降伏強度ＹＳ及び引張強度ＴＳの差が、それぞれＹＳで９０ＭＰａ以下かつＴＳで５０ＭＰａ以下であることを特徴とする請求項１又は２記載の衝撃特性に優れた歪時効硬化型熱延鋼構造部材。When the heat treatment is performed, the difference between the yield strength YS and the tensile strength TS of the portion deformed by the forming process and the undeformed portion is 90 MPa or less in YS and 50 MPa or less in TS, respectively. A strain age hardening type hot rolled steel structural member having excellent impact characteristics according to claim 1 or 2. 前記熱処理を行った際に、未変形部フェライト粒内における炭素・窒素原子集合体および炭窒化物析出物の２種合計の分布密度が、体積密度で１００〜１００００個・μｍ^−３であることを特徴とする請求項１〜３の何れか１　項に記載の衝撃特性に優れた歪時効硬化型熱延鋼構造部材。When the heat treatment is performed, the total distribution density of the two types of carbon / nitrogen atom aggregates and carbonitride precipitates in the undeformed portion ferrite grains is 100 to 10000 μm ⁻³ in volume density. The strain-age hardening type hot-rolled steel structural member excellent in impact characteristics according to any one of claims 1 to 3, characterized in that: 更に、質量％で、Ｎｉ：０．１〜２．０％を含むことを特徴とする請求項１〜４の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼構造部材。The strain-age-hardened hot-rolled steel structural member having excellent impact characteristics according to any one of claims 1 to 4, further comprising Ni: 0.1 to 2.0% by mass%. . 前記組成に加えてさらに、質量％で、下記ａ群〜ｄ群の１群または２群以上を含むことを特徴とする請求項１〜５の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼構造部材。
ａ群：Ｃｒ、Ｍｏ、Ｗのうち１種または２種以上の合計を０．１〜２．０％。
ｂ群：Ｎｂ、Ｔｉ、Ｖ、Ｔａのうち１種または２種以上の合計を０．００３〜
０．２％。
ｃ群：Ｂを０．０００３〜０．０１０％。
ｄ群：ＣａまたはＭｇのうち１種または２種を合計で０．００１〜０．０１％。The strain excellent in impact characteristics according to any one of claims 1 to 5, further comprising one or two or more groups of the following groups a to d in mass% in addition to the composition. Age hardening type hot rolled steel structural member.
Group a: 0.1 to 2.0% of the total of one or more of Cr, Mo, and W.
Group b: Nb, Ti, V, Ta, one or more of the total of 0.003 to
0.2%.
Group c: B 0.0003 to 0.010%.
Group d: 0.001 to 0.01% of one or two of Ca and Mg in total. 請求項１〜６の何れか１項に記載の熱延鋼材に電気めっき又は溶融めっきが施されていることを特徴とする衝撃特性に優れた歪時効硬化型熱延めっき鋼構造部材。A strain age hardening type hot rolled steel structural member excellent in impact characteristics, characterized in that the hot rolled steel material according to any one of claims 1 to 6 has been subjected to electroplating or hot dip plating. 請求項１、５又は６に記載の組成からなる鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　温度〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせによりＡ１　温度〜１００℃間を平均冷却速度０．０１〜３０℃／ｓで冷却することを特徴とする請求項１〜６の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼材を製造する方法。The steel slab having the composition according to claim 1, 5 or 6 is heated to a slab heating temperature of 1000 to 1270 ° C., rough-rolled, and then finish-rolled at an A1 ° temperature to 1000 ° C., air-cooled, water-cooled or The impact characteristics according to any one of claims 1 to 6, wherein cooling is performed at a mean cooling rate of 0.01 to 30 ° C / s between A1 ° and 100 ° C by a combination of these two cooling methods. For producing a strain-aged hardened hot-rolled steel with excellent heat resistance. 請求項１、５又は６記載の組成からなる鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　温度〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせによる冷却の後、Ａ１　温度〜１００℃間で巻取ることを特徴とする請求項１〜６の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼材を製造する方法。After heating the steel slab having the composition according to claim 1, 5 or 6 to a slab heating temperature of 1000 to 1270 ° C and performing rough rolling, finish rolling at an A1 ° temperature to 1000 ° C is completed, and air-cooling, water-cooling, or the like. The strain-age-hardened hot-rolled roll having excellent impact properties according to any one of claims 1 to 6, wherein after cooling by a combination of two kinds of cooling methods, winding is performed at a temperature of A1 ° to 100 ° C. A method of manufacturing steel. 請求項１、５又は６記載の組成を有する鋼スラブを、スラブ加熱温度１０００〜１２７０℃に加熱し、粗圧延をした後、Ａ１　温度〜１０００℃で仕上げ圧延を終了し、空冷、水冷あるいはこれら２種の冷却方法の組み合わせにより室温まで冷却したのち、１００℃〜Ａ１　温度の間で再加熱処理を行うことを特徴とする請求項１〜６の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼材を製造する方法。The steel slab having the composition according to claim 1, 5 or 6, is heated to a slab heating temperature of 1000 to 1270 ° C., rough-rolled, and then finish-rolled at an A1 ° temperature to 1000 ° C., air-cooled, water-cooled or The composition according to any one of claims 1 to 6, characterized in that after cooling to room temperature by a combination of two kinds of cooling methods, a reheating treatment is performed at a temperature of 100 ° C to A1 ° C. A method for producing a strain age hardening type hot rolled steel material. 請求項８〜１０の何れか１項に記載の製造工程に続いて、伸び率：０．５〜４０％の調質圧延またはレベラー加工を施すことを特徴とする請求項１〜６の何れか１項に記載の衝撃特性に優れた歪時効硬化型熱延鋼材を製造する方法。The tempering rolling or elevating at an elongation percentage of 0.5 to 40% is performed following the production process according to any one of claims 8 to 10. 2. A method for producing a strain-age hardened hot-rolled steel excellent in impact characteristics according to item 1.

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