JP2005036272A

JP2005036272A - Strain age hardening type steel excellent in cold non-aging property and burring workability, and its production method

Info

Publication number: JP2005036272A
Application number: JP2003199410A
Authority: JP
Inventors: Naoki Maruyama; 直紀丸山; Takehide Senuma; 武秀瀬沼; Masaaki Sugiyama; 昌章杉山
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 2003-07-18
Filing date: 2003-07-18
Publication date: 2005-02-10

Abstract

<P>PROBLEM TO BE SOLVED: To provide strain age hardening type steel satisfying a BH (Bake Hardening) quantity of 40 to 85 MPa, and jointly having cold non-aging properties, and to provide its production method. <P>SOLUTION: The strain age hardening type steel excellent in cold non-aging properties and burring workability has a composition comprising, by mass, 0.002 to 0.008% C, ≤0.5% Si, 0.1 to 3.0% Mn, ≤0.1% P, ≤0.1% S, 0.6 to 3.0% Cu, 0.1 to 3.0% Ni, ≥0.01% Al and ≤0.01% N, and in which N<SP>*</SP>specified by the following formula is ≤0, and the balance Fe with inevitable impurities. The area of Cu precipitates and Fe boundaries in the steel is ≥1[μm<SP>2</SP>/μm<SP>3</SP>], the area ratio of a ferritic phase is ≥85%, and the variation in yield point elongation ΔYPE1 by aging at 100°C for 1 hr is ≤0.6%: N<SP>*</SP>=N-14/27×Al. <P>COPYRIGHT: (C)2005,JPO&NCIPI

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常温保持中の品質劣化の少ない歪時効硬化型鋼材およびその製造方法に関するものであり、引張強度で３００ＭＰａから８００ＭＰａ程度の幅広い強度の鋼材に適用が可能である。特に、本発明鋼は優れたバーリング特性（伸びフランジ特性）を併せ持つことから、自動車の構造部材・足廻り部材として最適である。なお、本発明は塗装焼付処理工程を経て使用される建築物等の構造材料にも適用することが可能である。
【０００２】
【従来の技術】
車体重量軽減のニーズから自動車用鋼材においては高強度化の要請が高い。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形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。
【０００３】
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、成形加工時には軟質に保たれ、成形加工後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（ＢａｋｅＨａｒｄｅｎａｂｉｌｉｔｙ：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Ｃ原子あるいはＮ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Ｃ原子あるいはＮ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。
【０００４】
しかしながら、高いＢＨ量を得るために鋼材の固溶Ｃ量あるいは固溶Ｎ量を高めると常温時効劣化が生じ、その結果、加工時にストレッチャーストレインと呼ばれる歪み模様が発生したり、低加工歪み領域のＢＨ量が得られなくなるという問題があった。このように歪み時効硬化性と常温非時効特性の確保は二律相反するものと考えられており、実際、常温非時効性を確保した上で得られる最大のＢＨ量は高々３０〜４０ＭＰａ程度であった。
【０００５】
これを解決する手段として、特許文献１には焼鈍後の組織をフェライト相と低温変態生成相との複合組織とし、高ｒ値、高ＢＨ、高延性および常温非時効性を兼ね備えた冷延鋼板が開示されている。しかしながら、この技術には複合組織を得るためには極めて高い温度の焼鈍が必須となり、連続焼鈍時に板破断等のトラブルの原因となるという実操業上の問題点を有する。
【０００６】
また特許文献２には、Ｎｂを添加した極低炭素冷延鋼板において焼鈍後の冷却速度を制御することによって粒界中のＣ濃度を高めて、高ＢＨと常温遅時効性との両立が可能であることが示されている。しかしながら、常温非時効性を確保した上で得られるＢＨ量は高々５０ＭＰａ程度であり、また５５０ＭＰａ以上の高張力鋼に適用することも難しかった。
【０００７】
また特許文献３には、フェライトの結晶粒界中のＮ濃度を所定の範囲内に定めて、高ＢＨと常温遅時効性との両立を可能にする技術が示されている。しかしながら、この方法では常温時効時の全伸びの劣化は抑制されるものの、ストレッチャーストレイン発生の原因となる降伏点伸び発生の抑制には配慮が無い。さらにこの方法によると結晶粒が粗大になるほど常温遅時効を確保した上で得られるＢＨ量は減少するため、実鋼板への適用範囲は極めて限られるという問題点を有していた。
【０００８】
特許文献４にはＣｕ添加を特徴とする疲労特性に優れた加工用熱延鋼板に関する技術が開示されている。しかしながら、歪み時効硬化特性および常温非時効性に関する配慮はなされておらず、これらを制御するための最適条件は提示されていない。
【０００９】
【特許文献１】
特２８１８３１９号公報
【特許文献２】
特開平７−３００６２３号公報
【特許文献３】
特開２０００−２９７３５０号公報
【特許文献４】
特開平１１−２７９６９４号公報
【００１０】
【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、常温非時効性を兼備したＢＨ量４０〜８５ＭＰａの歪み時効硬化型鋼材およびその製造方法を提供することを目的とする。
【００１１】
【課題を解決するための手段】
本発明者らは上記の課題を達成するために、市中で最も広範に利用されている固溶Ｃ利用型のアルミキルド鋼材および単純ＳＵＬＣ（極低炭素）鋼材を中心に、高ＢＨを得た上で常温時効中での降伏点伸びの発生を抑制するための方法について検討を重ねた。その結果、鋼中に微細なＣｕ粒子を高密度で析出させ、さらにＣｕ析出相とＦｅの界面にＣを偏析させることにより、セメンタイトが析出し高いＢＨが得られなかったような緩冷却型の従来製造プロセスにおいても、常温非時効性を確保した上で高ＢＨを達成できることを見出した。すなわち、従来、固溶Ｃ利用型の歪み時効硬化型鋼材においては、高いＢＨ量を得るためには、熱延後あるいは焼鈍後の冷却中におけるセメンタイト形成を避けるために急冷のプロセスが必須であることが既に良く知られているが、Ｃｕ粒子を微細分散させておくとセメンタイトの形成が顕著に抑制され、その結果、急冷却のプロセスでなくても高いＢＨが得られ、さらに常温非時効性が得られることを明らかにした。
【００１２】
次いで、発明者らはＣ原子の存在状態に着目してこれら原因を調査した。その結果、セメンタイトの析出が抑制される原因は、Ｃｕ粒子とＦｅの界面（以下、Ｃｕ粒子／Ｆｅ界面ともいう）という偏析サイトが多量にあるために一カ所にＣ原子が集まりにくく、そのためセメンタイト析出が起こりにくくなっているためであり、一方、常温非時効と高ＢＨが両立される理由としては、１００℃以下の低温保持中においてはＣ原子はＣｕ析出物とＦｅの界面にトラップされて拡散しにくくなる一方で、１７０℃程度の高温では界面から脱離し、このために高ＢＨが得られているという全く新しいメカニズムによるものであることを見出し、本発明に至った。
【００１３】
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、常温非時効性とバーリング性に優れた歪時効硬化型鋼材であって、質量％で、
Ｃ：０．００２〜０．００８％、Ｓｉ：０．５％以下、
Ｍｎ：０．１〜３．０％、Ｐ：０．１％以下、
Ｓ：０．１％以下、Ｃｕ：０．６〜３．０％、
Ｎｉ：０．１〜３．０％、Ａｌ：０．０１％以上、
Ｎ：０．０１％以下
を含み、下式で規定されるＮ^＊が０以下であり、残部がＦｅおよび不可避的不純物からなり、鋼中におけるＣｕ析出物とＦｅの界面の面積が単位体積あたり１［μｍ^２／μｍ^３］以上であり、フェライトを面積率最大の相とし、１００℃で１時間の時効による降伏点伸びの変化量ΔＹＰ−Ｅｌが０．６％以下であることを特徴とする。
Ｎ^＊＝Ｎ−１４／２７×Ａｌ
【００１４】
（２）第２の発明は、前記組成に加えて、下記ａ群〜ｄ群の１群または２群以上を含有することを特徴とする。
ａ群：Ｃｒ，Ｍｏ，Ｗのうち１種または２種以上の合計を０．００５〜１．０％。
ｂ群：Ｎｂ，Ｔｉ，Ｖ，Ｔａのうち１種または２種以上の合計を０．００１〜０．２％。
ｃ群：Ｂを０．０００３〜０．０１０％。
ｄ群：Ｃａ，Ｍｇ，Ｚｒ，Ｃｅ，ＲＥＭのうち１種または２種以上を合計で０．００１〜０．０１％。
【００１５】
（３）第３の発明は、Ｃｕ析出物とＦｅの界面へのＮの偏析量ｎ_{ｅｘｃｅｓｓ}がＣｕ析出物とＦｅの界面の単位面積当たり０．０５［ａｔｏｍｓ／ｎｍ^２］以上であることを特徴とする前記（１）または（２）記載の歪み時効硬化型鋼板である。
【００１６】
（４）第４の発明は、前記（１）〜（３）の何れか１項に記載の歪み時効硬化型鋼板であって、電気めっきまたは溶融めっきが施されていることを特徴とする。
【００１７】
更に、本発明は常温非時効性に優れた歪時効硬化型鋼板の製造方法であって、（５）第５の発明は、前記（１）または（２）に記載の組成からなるスラブを加熱し、圧延終了温度（Ａｒ３ −１００）℃以上である熱間圧延を行い、５００〜６７０℃間で巻取りを行い、次いで１００〜３００℃間の滞留時間が２０ｓ以上である時効処理または冷却を行うことを特徴とする。
【００１８】
（６）第６の発明は、前記（１）または（２）に記載の組成からなるスラブを加熱し、圧延終了温度（Ａｒ３ −１００）℃以上である熱間圧延を行い、５５０〜７２０℃間の滞留時間が１７０ｓ以上となる冷却を行い、次いで１００〜３００℃間の滞留時間が２０ｓ以上である時効処理または冷却を行うことを特徴とする。
【００１９】
（７）第７の発明は、前記（１）または（２）に記載の組成からなる熱間圧延鋼材を冷間圧延した後、該冷延板をＡｃ１〜（Ａｃ１＋２００）℃間の最高到達温度で一次熱処理し、次いで５５０〜７２０℃間での滞留時間が１７０ｓ以上の時効析出処理を行い、さらに１００〜３００℃間の滞留時間が２０ｓ以上である時効処理または冷却を行うことを特徴とする。
【００２０】
（８）第８の発明は、前記（７）に記載の製造方法において、一次熱処理および時効析出処理の後、過時効処理を行い、次いで１００〜３００℃間の滞留時間が２０ｓ以上である時効処理または冷却を行うことを特徴とする。
【００２１】
（９）第９の発明は、前記（７）に記載の製造方法において、一次熱処理および時効析出処理を行った後に溶融亜鉛メッキ層を形成する工程を行うことを特徴とする。
【００２２】
（１０）第１０の発明は、前記（９）に記載の製造方法において、溶融亜鉛めっき層を形成した後、合金化処理を行うことを特徴とする。
【００２３】
（１１）第１１の発明は、前記（５）〜（１０）の何れか１項に記載の方法により製造した鋼板に、伸び率：３％以下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。
【００２４】
【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、０．００２％未満であると４０ＭＰａを超えるＢＨ量が得られず、一方、０．００８％を超えるとセメンタイト析出が起こり、常温で非時効化することが難しくなる。このため、本発明ではＣの範囲を０．００２〜０．００８％に限定した。
【００２５】
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。しかしながら、０．５％を超えると化成処理性やめっきの密着性が悪くなる。従ってＳｉ含有量を０．５％以下の範囲に制限した。０．３％以下がより好ましい範囲である。下限は特に限定することなく本発明の効果を奏することができるが、不純物として不可避的に０．００１％以上含有する場合が多い。
【００２６】
Ｍｎ：Ｍｎ：Ｍｎは鋼材のミクロ組織および強度の調整に用いられ、さらにＣｕ粒子の析出を促進させる効果がある。Ｍｎの含有量が０．１％未満であるとＣｕ粒子の析出に時間がかかるため生産性が急激に低下し、一方、３．０％を超えると成形加工性の劣化を招く。従って、Ｍｎ含有量を０．１〜３．０％の範囲に制限した。なお、Ｃｕ粒子の析出処理時間をさらに短縮するという観点から０．３％以上の添加が望ましい。
【００２７】
Ｐ：Ｐは熱延組織の微細化能を有し、また強力な固溶強化元素であることから鋼材の強度の調整に用いられる。ただし、添加量が０．１％を超えると、スポット溶接後の疲労強度が劣悪となったり、降伏強度が増加し過ぎてプレス時に面形状不良を引き起こす。さらに、連続溶融亜鉛メッキ時に合金化反応が極めて遅くなり、生産性が低下する。また、２次加工性も劣化する。従ってＰ含有量の範囲を０．１％以下に制限した。下限は特に限定することなく本発明の効果を奏することができるが、不純物として不可避的に０．００１％以上含有する場合が多い。
【００２８】
Ｓ：ＳはＭｎＳ，ＣｕＳ，ＴｉＳ等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。しかしながら、０．１％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０．１％以下に限定した。下限は特に限定することなく本発明の効果を奏することができるが、不純物として不可避的に０．０００１％以上含有する場合が多い。
【００２９】
Ｃｕ：Ｃｕは本発明において必須の添加元素である。０．６％未満であるとフェライト中にＣｕ粒子が析出せず、また３．０％を超えると熱間加工割れが起こり、またコスト的にも割高になる。従って、その適正添加範囲を０．６％〜３．０％に限定した。
【００３０】
Ｎｉ：ＮｉはＣｕ添加による熱間加工割れを抑制するために用いられる。０．１％未満であると熱間割れを抑制することが難しく、一方、３．０％を超えるとコスト的に割高となる。従って、その適正添加範囲を０．１％〜３．０％に限定した。
【００３１】
Ａｌ：Ａｌは鋼材のミクロ組織および強度の調整および脱酸調製元素として使用する。Ａｌ量が０．０１％未満であると固溶Ｎが残留する可能性がある。従って、Ａｌ量の範囲０．０１％以上に制限した。
【００３２】
Ｎ：Ｎは窒化物として、主にオーステナイト域の結晶粒径制御に用いられる。一方、Ｎが０．０１％を超えると、フェライト粒内に多量の炭窒化物が析出しＣｕの添加効果が弱まるため、Ｎ含有量の範囲を０．０１％以下とした。
【００３３】
Ｎ^＊：有効固溶Ｎ量Ｎ^＊が０を超えるとＮがＢＨに寄与するので、ＢＨ量の制御が難しくなる。このため本発明ではＮ^＊の範囲を０以下に限定した。
【００３４】
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０．００５〜１．０％。Ｃｒ，Ｍｏ，Ｗは炭窒化物形成元素であり、これらの元素の１種または２種以上の合計を０．００５％以上含有することにより、熱間圧延中、冷却中、あるいは一次熱処理工程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。しかしながら、１種または２種以上の合計で１．０％を超えると炭窒化物の析出量が多くなり、高ＢＨを得ることが困難になる場合があり、また成形加工性の劣化も招く。従って、その合計量の範囲を０．００５〜１．０％とした。
【００３５】
ｂ群：Ｎｂ，Ｔｉ，Ｖ，Ｔａのうち１種または２種以上の合計を０．００１〜０．２％。
Ｎｂ，Ｔｉ，Ｖ，Ｔａは炭窒化物形成元素であり、鋼材のミクロ組織、集合組織およびＣ量、Ｎ量を調整するのに用いられるので、１種または２種以上の合計を０．００１％以上含有することが好ましい。しかしながら、１種または２種以上の合計で０．２％を超えると、炭窒化物の析出量が多くなり、高ＢＨを得ることが困難になり、０．００１％未満では添加効果が現れない。従って、その合計量の範囲を０．００１〜０．２％とした。
【００３６】
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。
【００３７】
ｄ群：Ｃａ，Ｍｇ，Ｚｒ，Ｃｅ，ＲＥＭのうち１種または２種以上を合計で０．００１〜０．０１％。
Ｃａ，Ｍｇ，Ｚｒ，ＣｅおよびＲＥＭは介在物の形態、分布の制御に用いる元素であり、１種または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、成形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。なお、本発明において、ＲＥＭとはＬａおよびランタノイド系列の元素を指すものとする。
【００３８】
なお、不可避不純物として重要な元素としてＯがある。Ｏ量は脱酸の方法によりその残留量が大きく変化するが、不可避的に０．０００５％以上含有する場合が多い。
【００３９】
本発明に係る鋼材はＣｕ析出物／Ｆｅ界面でのＣ原子トラップ効果を十分に発現させ、さらに優れた疲労特性と穴拡げ性を得るために、フェライト単相組織であることが好ましく、少なくとも面積率最大の相がフェライトである必要がある。なお、本発明において「フェライト」とは、ＩＳＩＪｉｎｔｅｒｎａｔｉｏｎａｌ３５巻（１９９５）９４１〜９４４頁に示すようなポリゴナルフェライト、擬ポリゴナルフェライトあるいはグラニュラーベイニティックフェライトを指す。フェライト以外の残部組織はマルテンサイト、オーステナイト、パーライトの１種または２種以上を合計で２％まで含有しても良い。なお、フェライトの平均結晶粒径は特に限定する必要はなく、あらゆる結晶粒径で本発明の効果を奏功することができる。ただし、Ｃｕ粒子の粒界析出を防止する観点から３μｍ以上の粒径であることが望ましい。なお、本発明鋼板は軟質のフェライト組織を主体にしてＣｕ粒子の析出により高強度化することを特徴とするので、複相組織化により高強度化する鋼板に比べて格段に優れたバーリング特性を有する。
【００４０】
鋼中における単位体積あたりのＣｕ析出物／Ｆｅ界面の面積は、１［μｍ^２／μｍ^３］未満であるとＮ原子を十分にトラップすることができず、その結果、常温非時効性が得られない。従って、その界面面積量の範囲を１［μｍ^２／μｍ^３］以上に制限した。１０以上がより好ましい範囲である。
【００４１】
またＣｕ析出物／Ｆｅ界面へのＣの単位界面面積あたりの偏析量が合計で０．０５［ａｔｏｍｓ／ｎｍ^２］未満であると常温非時効性が得られない。従って、界面偏析量の範囲を０．０５［ａｔｏｍｓ／ｎｍ^２］以上に制限した。より安定した常温非時効性を得るためには０．２以上、より好ましくは１以上であり、Ｃｕ粒子を成長させＣｕ析出物／Ｆｅ界面を非整合界面にすることでより大きな偏析量を実現することができる。
【００４２】
常温非時効性は人工時効後の降伏点伸びによって評価するのが簡易で好適である。本発明によって得られる鋼材は、１００℃にて１時間熱処理後の引張試験による降伏点伸びの増加量が０．６％以下である。なお、本発明鋼材は常温時効により降伏点伸び発生が抑制されるだけでなく、全伸び値の低下も抑制される。この原因は粒界部における鉄炭窒化物の生成が抑制されているためであると推測され、ＮやＣが鋼中のＣｕ粒子／Ｆｅ界面にトラップされることに起因する。具体的には全伸びの劣化量は２％以下に抑制される。
【００４３】
次に、製造方法の限定理由について説明する。
熱間圧延に供するスラブは特に限定するものではない。すなわち、連続鋳造スラブや薄スラブキャスターなどで製造したものであればよい。また、鋳造後に直ちに熱間圧延を行う連続鋳造−直接圧延（ＣＣ−ＤＲ）のようなプロセスにも適合する。
【００４４】
熱延の条件については、巻取中あるいは冷却中においてマトリクス中にＣｕ粒子をより微細に析出させる観点から、（Ａｒ３ −１００）℃以上の仕上げ終了温度にする必要がある。（Ａｒ３ −１００）℃未満では、Ｃｕ粒子の転位上析出によるＣｕ粒子の粗大析出化を引き起こしたり、また板厚精度の問題を生じたりする。Ａｒ３温度以上がより好ましい範囲である。仕上げ温度の上限は特に定めることなく本発明の効果を得ることができるが、ｒ値を確保するためには１０００℃以下とすることが好ましい。
【００４５】
熱延後はＣｕをＦｅマトリクス中に析出させるために、５００℃〜６７０℃間で巻取処理を行うか（前記（５）に係る本発明）、あるいは５５０〜７２０℃間の滞留時間が１７０ｓ以上になるように冷却する必要がある（前記（６）に係る本発明）。巻取温度が６７０℃を超えるとＣｕ粒子が粗大化して、１［μｍ^２／μｍ^３］以上のＣｕ粒子／Ｆｅ界面面積を得ることができず、その結果、常温非時効性が得られない。巻取温度が５００℃未満であると通常の巻取時間内にＣｕ粒子の析出が起きない。また、５５０〜７２０℃間の滞留時間が１７０ｓ未満であるとＣｕの析出が起こらない。
【００４６】
上記の冷却あるいは巻取後は１００〜３００℃間で２０ｓ以上滞留させる。詳細なメカニズムは定かではないが、この工程中にＣｕ粒子／Ｆｅ界面への固溶Ｃの偏析が起こると考えられ、常温保持中の降伏点伸び発現抑制を得るために必須の工程である。１００〜３００℃間が結晶粒内において以上の原子移動を最も迅速に起こさせるための温度域であり、この温度範囲内の滞留時間が２０ｓ未満ではＢＨ性と常温遅時効性の両立することができない。従って、滞留時間の範囲を２０ｓ以上に制限した。なお、より優れた常温遅時効性を得るという観点からは、６０ｓ以上の保持がより好ましい。滞留時間の上限については特に定めないが、６００００ｓを超えると粒界への粗大炭窒化物の析出が起こり、高ＢＨが得られない場合があり、また伸び値が顕著に低下するので、その上限としては６００００ｓ以下であることが好ましい。
【００４７】
熱延後は必要に応じて酸洗し、その後インラインまたはオフラインで圧下率３％以下の調質圧延若しくはレベラー加工、または圧下率４０％程度までの冷間圧延を施しても構わない（前記（１１）に係る本発明）。
【００４８】
次に冷延板あるいはめっき板を最終製品とする場合の製造条件について示す。
素材である熱延板の製造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。続いて酸洗等の通常公知の処理を行い、冷間圧延を行う。冷間圧延の条件については、圧延パスの回数、圧下率については特に規定する必要はなく常法に従えばよい。ただし、冷間圧延の圧下率が９０％超では設備への負荷が過大となり、さらに製品の機械的性質の異方性が大きくなるので、９０％以下であることが好ましい。
【００４９】
連続焼鈍工程または連続焼鈍及びめっき工程における加熱速度については常法に従えばよい。一次熱処理時の最高到達温度については、Ａｃ１温度未満では再結晶が完了せず加工性が劣悪となったり、セメンタイトが十分に固溶せずＢＨが得られない場合がある。一方、一次熱処理温度が（Ａｃ１＋２００）℃を超えると、通板中に板破断をおこし生産性が大きく低下する。従って一次熱処理の最高到達温度の範囲をＡｃ１〜（Ａｃ１＋２００）℃に制限した。
【００５０】
上記一次熱処理終了後、ＣｕをＦｅマトリクス中に析出させるために、５５０〜７２０℃間の滞留時間が１７０ｓ以上の析出処理を行う必要がある。５５０〜７２０℃間の滞留時間が１７０ｓ未満である場合にはＣｕの析出が起こらず、常温非時効性の確保が困難になる（前記（７）に係る本発明）。
【００５１】
溶融亜鉛めっきを施さない場合は、一次熱処理後、上記の５００℃までの冷却をした後は、過時効処理として５００℃〜室温までの何れの温度まで冷却しても構わない（前記（８）に係る本発明）。
【００５２】
溶融亜鉛めっきを施す場合には、一次熱処理後、少なくとも５５０〜７２０℃間の滞留時間が１７０ｓ以上の析出処理を行い、亜鉛めっきを行い、その後必要に応じてめっき相の合金化処理を行う。亜鉛めっきおよび合金化の条件は特に定めないが、添加したＣの粒界への析出を抑止する観点からめっき浴中への浸漬時間および合金化炉中の保持時間はそれぞれ４０ｓ以下、より好ましくは２０ｓ以下であることが好ましい（前記（９），（１０）に係る本発明）。
【００５３】
一次熱処理後、あるいは過時効帯を利用した熱処理後、あるいはめっき処理後（合金化処理後も含む）に１００〜３００℃間の滞留時間が２０ｓ以上である冷却を行う。なお、過時効処理を行なう場合には、前述の温度域で滞留時間を確保するため、１００〜３００℃、２０ｓ以上とすることが熱効率上好ましい。なお、１００〜３００℃間の滞留時間が２０ｓ以上である冷却工程は上に述べたように常温保持中の降伏点伸び発現抑制を得るために必須の工程である。１００〜３００℃間は結晶粒内において以上の原子移動を最も迅速に起こさせるための温度域であり、この温度範囲内の滞留時間が２０ｓ未満ではＢＨ性と常温非時効性を両立することができない。従って、滞留時間の範囲を２０ｓ以上に制限した。なお、より優れた常温非時効性を得るという観点からは、６０ｓ以上の保持がより好ましい。滞留時間の上限については特に定めないが、６００００ｓを超えると粒界への粗大炭窒化物の析出が起こり、高ＢＨが得られない場合があり、また伸び値が顕著に低下するので、その上限としては６００００ｓ以下であることが好ましい。
【００５４】
冷延板またはめっき板を最終製品とする場合も調質圧延またはレベラー加工は、常温遅時効性の向上と形状矯正のために行い、圧下率３％以下の範囲で行うのがよい。３％を超えるとＢＨ量が低下する傾向があるので、これを上限とする（前記（１１）に係る本発明）。
【００５５】
１次熱処理後にめっき工程あるいはめっき合金化工程を経ずに作られた本発明の冷延鋼板は、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アルミ、亜鉛、クロム、錫、ニッケルが例として挙げられる。
【００５６】
なお、Ｃｕ粒子／Ｆｅ間の界面面積は透過電子顕微鏡により像観察を行い決定するか、あるいはＣｕ析出相が極めて微細な場合はアトムプローブ電界イオン顕微鏡法により測定領域内の平均Ｃｕ粒子径および析出物の体積分率を測定し決定する方法が好適である。Ｃｕ粒子／Ｆｅ界面へのＮあるいはＣの偏析量はアトムプローブ電界イオン顕微鏡法により測定する方法が簡易で好適である。なおＣｕ粒子のサイズについては直径１ｎｍ以上のものを析出物と見なすことにする。
【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に示す成分の鋼を溶製し、表２に示す条件で熱間圧延工程を行った。なお、表２に示すように熱間圧延時のスラブ加熱温度は１０５０〜１２５０℃で、最終板厚を４ｍｍとした。調質圧延は全て１．０％の伸び率で行った。このようにして得られた鋼板について、引張試験、ＢＨ試験および組織観察を行った。
【００５８】
また、表１に示す成分の鋼について、１０５０〜１２５０℃にスラブを再加熱し、熱間圧延終了温度８４０〜９３０℃で最終板厚４ｍｍまで熱延し、５５０〜６５０℃で巻取り、このようにして得られた熱延鋼板を酸洗の後、７０〜８５％の冷延率で冷間加工を行い、脱脂処理を行ったのち、表３に示す条件で連続熱処理および連続亜鉛めっき工程を行った。調質圧延率は全て１．０％で行った。このようにして得られた鋼板について、引張試験、ＢＨ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。
【００５９】
引張試験はＪＩＳ５号試験片を用い、歪み速度１０^−３／ｓの条件で行った。常温保持中の材質変化は、１００℃×１ｈｒの促進時効前後の引張試験結果を比較することにより評価した。一方、ＢＨ量を観察するための引張試験はＪＩＳ１３Ｂ試験片を用い、歪み速度１０^−３／ｓの条件で行った。ＢＨ試験の予変形量は２％、塗装焼付処理に対応する時効条件は１７０℃×２０分で行い、再引張時において上部降伏点で評価したＢＨ量をとった。フェライトの平均結晶粒径はＪＩＳＧ０５５２の試験方法に従って行った。試験結果を表４に示す。
バーリング加工性（伸びフランジ性）については日本鉄鋼連盟規格ＪＦＳＴ
１００１−１９９６記載の穴拡げ試験方法によって評価した。
【００６０】
試料Ｎｏ．１，Ｎｏ．８はＣｕ量が適正範囲外で十分な量のＣｕ粒子／Ｆｅ界面面積が得られず、その結果常温非時効特性が得られなかった例である。試料Ｎｏ．４，Ｎｏ．１２は熱延巻取温度あるいは７２０℃〜５５０℃間の滞留時間が適正範囲外であるために、十分な量のＣｕ粒子／Ｆｅ界面面積が得られず、その結果常温非時効特性が得られなかった例である。試料Ｎｏ．５，Ｎｏ．１０は、３００〜１００℃間の滞留時間が適正範囲外であったためにＣｕ粒子／Ｆｅ界面への十分な量のＣ偏析が得られず、その結果常温非時効特性が得られなかった例である。試料Ｎｏ．７，Ｎｏ．１４はＡｌ量に対してＮ量が過剰になり、ＢＨ量が過大になったため、降伏点伸びの変化量ΔＹＰ−ＥＬが０．６％以内に収まらなかった例である。
【００６１】
【表１】

【００６２】
【表２】

【００６３】
【表３】

【００６４】
【表４】

【００６５】
【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、成形限界値が優れ、常温保持中の材質劣化が少なく、バーリング加工性に優れ、高い歪み硬化能を有する歪み時効硬化型鋼材を安価に提供することができ、工業的に価値が高い。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a strain age-hardening type steel material with less quality deterioration during normal temperature holding, which is suitable for the use of structural members, suspension members and panel members of automobiles, and a method for producing the same, and a tensile strength of about 300 MPa to 800 MPa. It can be applied to a wide range of strength steel materials. In particular, the steel according to the present invention has an excellent burring characteristic (stretch flange characteristic), and is therefore optimal as a structural member / suspension member for automobiles. In addition, this invention is applicable also to structural materials, such as a building used through a paint baking process.
[0002]
[Prior art]
Due to the need to reduce vehicle weight, there is a strong demand for higher strength in automotive steel. However, it is generally known that increasing the strength of a material is accompanied by deterioration of press formability such as a decrease in shape freezing property and cracking during molding, and there is a strong method of increasing the strength without reducing workability. It was desired.
[0003]
In response to such demands, as a technology to achieve high strength while ensuring moldability, the strain age-hardening phenomenon that occurs during the electrodeposition coating baking process after molding is used. A technique using so-called bake hardenability (BH) that increases yield strength or tensile strength is known. This type of steel material has C atoms or N atoms dissolved during forming to ensure formability, and in the electrodeposition coating baking process, C atoms or N atoms are fixed to dislocations generated in the steel during forming. Or yield strength or tensile strength is increased by finely dispersing and depositing carbide or nitride on the dislocations.
[0004]
However, increasing the solute C content or solute N content of the steel material in order to obtain a high BH content will cause aging deterioration at room temperature, resulting in a strain pattern called stretcher strain during processing, There was a problem that the amount of BH could not be obtained. Thus, it is considered that securing the strain age-hardening property and the non-aging property at room temperature is contradictory, and the maximum amount of BH obtained after securing the non-aging property at room temperature is about 30 to 40 MPa at most. there were.
[0005]
As means for solving this, Patent Document 1 discloses a cold-rolled steel sheet having a composite structure of a ferrite phase and a low-temperature transformation phase as a structure after annealing and having a high r value, high BH, high ductility, and non-aging at room temperature. Is disclosed. However, this technique requires an extremely high temperature annealing in order to obtain a composite structure, and has a problem in actual operation that causes troubles such as plate breakage during continuous annealing.
[0006]
In Patent Document 2, it is possible to increase the C concentration in the grain boundary by controlling the cooling rate after annealing in the ultra-low carbon cold-rolled steel sheet to which Nb is added, and to achieve both high BH and room temperature slow aging. It is shown that. However, the amount of BH obtained after securing non-aging at room temperature is about 50 MPa at most, and it was difficult to apply it to high-tensile steel of 550 MPa or more.
[0007]
Further, Patent Document 3 discloses a technique that enables both high BH and room temperature slow aging by setting the N concentration in the crystal grain boundary of ferrite within a predetermined range. However, although this method suppresses the deterioration of the total elongation during aging at room temperature, there is no consideration for the suppression of the yield point elongation that causes the occurrence of stretcher strain. Furthermore, according to this method, as the crystal grains become coarser, the amount of BH obtained after securing the room temperature slow aging decreases, so that there is a problem that the range of application to actual steel sheets is extremely limited.
[0008]
Patent Document 4 discloses a technique related to a hot-rolled steel sheet for machining excellent in fatigue characteristics characterized by Cu addition. However, no consideration is given to strain age hardening characteristics and non-aging properties at room temperature, and no optimum condition for controlling these has been proposed.
[0009]
[Patent Document 1]
Japanese Patent No. 2818319 [Patent Document 2]
JP-A-7-300623 [Patent Document 3]
JP 2000-297350 A [Patent Document 4]
JP-A-11-279694 [0010]
[Problems to be solved by the invention]
The present invention has been made in view of the actual situation as described above, and is used for automotive structural members, suspension members, panel members, architectural structural members, and inner and outer plates of electrical products, which are made through an electrodeposition coating process. An object of the present invention is to provide a strain age-hardening type steel material having a BH amount of 40 to 85 MPa and a manufacturing method thereof suitable for a panel and having a non-aging property at room temperature.
[0011]
[Means for Solving the Problems]
In order to achieve the above-mentioned problems, the present inventors have obtained a high BH centering on a solid solution C-utilized aluminum killed steel material and a simple SULC (very low carbon) steel material that are most widely used in the city. The method for suppressing the occurrence of yield point elongation during aging at room temperature was studied repeatedly. As a result, a slow cooling type in which cementite is precipitated and high BH cannot be obtained by precipitating fine Cu particles in steel at high density and segregating C at the interface between the Cu precipitation phase and Fe. It has also been found that high BH can be achieved in a conventional manufacturing process while ensuring non-aging at room temperature. That is, conventionally, in a strain-age hardening type steel material using a solute C, in order to obtain a high BH amount, a rapid cooling process is indispensable in order to avoid formation of cementite during cooling after hot rolling or annealing. It is already well known, but when the Cu particles are finely dispersed, the formation of cementite is remarkably suppressed, and as a result, a high BH can be obtained even if it is not a rapid cooling process. It was clarified that can be obtained.
[0012]
Next, the inventors investigated these causes by paying attention to the existence state of C atoms. As a result, the precipitation of cementite is suppressed because there are a large number of segregation sites called the interface between Cu particles and Fe (hereinafter also referred to as Cu particles / Fe interface), making it difficult for C atoms to collect in one place. This is because precipitation is less likely to occur. On the other hand, the reason why both non-aging at normal temperature and high BH are compatible is that C atoms are trapped at the interface between Cu precipitates and Fe during low temperature holding at 100 ° C. or lower. While becoming difficult to diffuse, the inventors have found that this is due to a completely new mechanism of desorption from the interface at a high temperature of about 170 ° C., and thus high BH is obtained, leading to the present invention.
[0013]
The present invention has taken the following means in order to solve the above problems. That is,
(1) The first invention is a strain age-hardening steel material excellent in non-aging at normal temperature and burring properties, and is in mass%.
C: 0.002 to 0.008%, Si: 0.5% or less,
Mn: 0.1 to 3.0%, P: 0.1% or less,
S: 0.1% or less, Cu: 0.6-3.0%,
Ni: 0.1 to 3.0%, Al: 0.01% or more,
N: 0.01% or less, N ^* defined by the following formula is 0 or less, the balance is Fe and inevitable impurities, and the area of the interface between Cu precipitates and Fe in the steel is per unit volume. 1 [μm ² / μm ³ ] or more, with ferrite as the phase with the largest area ratio, and a change in yield point elongation ΔYP-El due to aging at 100 ° C. for 1 hour is 0.6% or less. To do.
N ^* = N-14 / 27 × Al
[0014]
(2) In addition to the said composition, 2nd invention contains 1 group or 2 groups or more of the following a group-d group, It is characterized by the above-mentioned.
Group a: 0.005 to 1.0% of the total of one or more of Cr, Mo, and W.
Group b: 0.001 to 0.2% of the total of one or more of Nb, Ti, V, and Ta.
c group: B is 0.0003 to 0.010%.
d group: One or more of Ca, Mg, Zr, Ce, and REM in total 0.001 to 0.01%.
[0015]
(3) In the third invention, the amount of N segregation n _{excess at} the interface between the Cu precipitate and Fe is 0.05 [atoms / nm ² ] or more per unit area of the interface between the Cu precipitate and Fe. The strain age-hardened steel sheet according to (1) or (2), which is characterized in that
[0016]
(4) A fourth invention is the strain age hardening type steel sheet according to any one of (1) to (3), wherein electroplating or hot dipping is performed.
[0017]
Furthermore, the present invention is a method for producing a strain age-hardening type steel sheet having excellent non-aging properties at room temperature, wherein (5) the fifth invention is for heating a slab comprising the composition described in (1) or (2) above. Then, hot rolling at a rolling end temperature (Ar3-100) ° C. or higher is performed, winding is performed at 500 to 670 ° C., and then an aging treatment or cooling in which a residence time between 100 to 300 ° C. is 20 seconds or longer. It is characterized by performing.
[0018]
(6) 6th invention heats the slab which consists of a composition as described in said (1) or (2), performs hot rolling which is more than rolling completion temperature (Ar3-100) degreeC, 550-720 degreeC. Cooling is performed such that the residence time is 170 s or more, and then an aging treatment or cooling in which the residence time between 100 to 300 ° C. is 20 s or more is performed.
[0019]
(7) In the seventh invention, after cold-rolling a hot-rolled steel material having the composition described in (1) or (2) above, the cold rolled sheet is made to reach the maximum between Ac1 and (Ac1 +200) ° C. It is characterized by performing a primary heat treatment at a temperature, then performing an aging precipitation treatment with a residence time between 550 and 720 ° C. of 170 s or more, and further performing an aging treatment or cooling with a residence time between 100 and 300 ° C. of 20 s or more. To do.
[0020]
(8) According to an eighth aspect of the present invention, in the production method according to (7), after the primary heat treatment and the aging precipitation treatment, an overaging treatment is performed, and then the residence time between 100 to 300 ° C. is 20 s or more. It is characterized by performing treatment or cooling.
[0021]
(9) A ninth invention is characterized in that, in the manufacturing method according to (7), a step of forming a hot-dip galvanized layer is performed after the primary heat treatment and the aging precipitation treatment.
[0022]
(10) The tenth invention is characterized in that, in the manufacturing method according to the above (9), an alloying treatment is performed after forming a hot-dip galvanized layer.
[0023]
(11) The eleventh invention is characterized by subjecting a steel sheet produced by the method according to any one of (5) to (10) to temper rolling or leveler processing with an elongation of 3% or less. And
[0024]
DETAILED DESCRIPTION OF THE INVENTION
The present invention is described in detail below.
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 of steel and the control of the microstructure. However, if it is less than 0.002%, a BH amount exceeding 40 MPa cannot be obtained. On the other hand, if it exceeds 0.008%, cementite precipitation occurs and it becomes difficult to non-age at room temperature. For this reason, in the present invention, the range of C is limited to 0.002 to 0.008%.
[0025]
Si: Si is used to adjust the microstructure and strength of steel. However, if it exceeds 0.5%, the chemical conversion treatment property and the adhesion of plating deteriorate. Therefore, the Si content is limited to a range of 0.5% or less. 0.3% or less is a more preferable range. The lower limit is not particularly limited, and the effects of the present invention can be achieved. However, the lower limit is inevitably contained as 0.001% or more in many cases.
[0026]
Mn: Mn: Mn is used for adjusting the microstructure and strength of the steel material, and further has an effect of promoting the precipitation of Cu particles. If the Mn content is less than 0.1%, it takes a long time to precipitate Cu particles, so that the productivity is drastically reduced. On the other hand, if it exceeds 3.0%, the moldability is deteriorated. Therefore, the Mn content is limited to a range of 0.1 to 3.0%. In addition, addition of 0.3% or more is desirable from the viewpoint of further shortening the Cu particle precipitation treatment time.
[0027]
P: P has the ability to refine a hot-rolled structure and is a strong solid solution strengthening element, so it is used for adjusting the strength of steel. However, if the addition amount exceeds 0.1%, the fatigue strength after spot welding becomes poor, or the yield strength increases excessively, causing surface shape defects during pressing. Furthermore, the alloying reaction becomes extremely slow during continuous hot dip galvanizing, and productivity is lowered. Also, the secondary workability is deteriorated. Therefore, the range of P content is limited to 0.1% or less. The lower limit is not particularly limited, and the effects of the present invention can be achieved. However, the lower limit is inevitably contained as 0.001% or more in many cases.
[0028]
S: S is present in steel as MnS, CuS, TiS, etc., and is used to adjust the strength and ductility of the steel material through control of the crystal grain size. However, if it exceeds 0.1%, hot brittleness may occur, so the range was limited to 0.1% or less. The lower limit is not particularly limited, and the effects of the present invention can be achieved. However, the lower limit is inevitably contained in an amount of 0.0001% or more.
[0029]
Cu: Cu is an essential additive element in the present invention. If it is less than 0.6%, Cu particles do not precipitate in the ferrite, and if it exceeds 3.0%, hot working cracks occur, and the cost is high. Therefore, the appropriate addition range is limited to 0.6% to 3.0%.
[0030]
Ni: Ni is used to suppress hot working cracks due to Cu addition. If it is less than 0.1%, it is difficult to suppress hot cracking, while if it exceeds 3.0%, the cost becomes high. Therefore, the appropriate addition range is limited to 0.1% to 3.0%.
[0031]
Al: Al is used as an element for adjusting the microstructure and strength of steel and for deoxidation. If the Al content is less than 0.01%, solid solution N may remain. Therefore, the Al content is limited to 0.01% or more.
[0032]
N: N is used as a nitride mainly for controlling the crystal grain size in the austenite region. On the other hand, if N exceeds 0.01%, a large amount of carbonitride precipitates in the ferrite grains and the effect of adding Cu is weakened, so the range of N content is set to 0.01% or less.
[0033]
N ^* : When the effective solid solution N amount N ^* exceeds 0, N contributes to BH, so that it becomes difficult to control the BH amount. Therefore, in the present invention, the range of N ^* is limited to 0 or less.
[0034]
In the present invention, in addition to the above-described composition, the object of the present invention can be achieved even if one or more of the groups a to d are contained.
Group a: The total of one or more of Cr, Mo, and W is 0.005 to 1.0%. Cr, Mo, and W are carbonitride-forming elements. By containing 0.005% or more of one or more of these elements in total, during hot rolling, during cooling, or during a primary heat treatment step It is used to adjust the strength of the steel material by precipitating mainly as carbonitride. However, if the total of one or more types exceeds 1.0%, the amount of carbonitride deposited increases, and it may be difficult to obtain high BH, and the moldability may be deteriorated. Therefore, the range of the total amount is set to 0.005 to 1.0%.
[0035]
Group b: 0.001 to 0.2% of the total 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, texture, C content, and N content of steel materials. % Or more is preferable. However, if the total of one or more types exceeds 0.2%, the amount of carbonitride deposited increases, making it difficult to obtain high BH. If it is less than 0.001%, the effect of addition does not appear. . Therefore, the range of the total amount is set to 0.001 to 0.2%.
[0036]
c group: B is 0.0003 to 0.010%.
When B is contained in an amount of 0.0003% or more, it segregates at the grain boundary, has the effect of suppressing secondary processing cracks due to P, and has the effect of improving the moldability. 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%.
[0037]
d group: One or more of Ca, Mg, Zr, Ce, and REM in total 0.001 to 0.01%.
Ca, Mg, Zr, Ce and REM are elements used for controlling the form and distribution of inclusions, and preferably contain one or two or more in total of 0.001% or more. However, if the total content exceeds 0.01%, the moldability is deteriorated. Therefore, the total amount range is set to 0.001 to 0.01%. In the present invention, REM refers to La and lanthanoid series elements.
[0038]
Note that O is an important element as an inevitable impurity. Although the amount of O varies greatly depending on the deoxidation method, it is unavoidably contained in an amount of 0.0005% or more.
[0039]
The steel material according to the present invention preferably has a ferrite single-phase structure in order to sufficiently exhibit the C atom trapping effect at the Cu precipitate / Fe interface and to obtain excellent fatigue characteristics and hole expansibility, and has at least an area. The phase with the highest rate needs to be ferrite. In the present invention, “ferrite” refers to polygonal ferrite, pseudo-polygonal ferrite, or granular bainitic ferrite as shown in ISIJ International 35 (1995) pages 941-944. The remaining structure other than ferrite may contain up to 2% in total of one or more of martensite, austenite, and pearlite. The average crystal grain size of ferrite is not particularly limited, and the effect of the present invention can be achieved with any crystal grain size. However, it is desirable that the particle size is 3 μm or more from the viewpoint of preventing the grain boundary precipitation of Cu particles. In addition, since the steel sheet of the present invention has a soft ferrite structure as a main component and is strengthened by precipitation of Cu particles, it has a burring characteristic that is remarkably superior to a steel sheet that has a high strength by forming a multiphase structure. Have.
[0040]
If the area of the Cu precipitate / Fe interface per unit volume in steel is less than 1 [μm ² / μm ³ ], N atoms cannot be sufficiently trapped, resulting in room temperature non-aging. I can't. Therefore, the range of the interface area amount is limited to 1 [μm ² / μm ³ ] or more. 10 or more is a more preferable range.
[0041]
Further, when the total amount of segregation per unit interface area of C to the Cu precipitate / Fe interface is less than 0.05 [atoms / nm ² ], the room temperature non-aging property cannot be obtained. Therefore, the range of the amount of interface segregation was limited to 0.05 [atoms / nm ² ] or more. In order to obtain a more stable normal temperature non-aging property, it is 0.2 or more, more preferably 1 or more, and a larger amount of segregation is realized by growing Cu particles and making the Cu precipitate / Fe interface non-matching interface. can do.
[0042]
It is simple and preferable to evaluate the non-aging property at room temperature by the yield point elongation after artificial aging. In the steel material obtained by the present invention, the increase in yield point elongation by a tensile test after heat treatment at 100 ° C. for 1 hour is 0.6% or less. The steel of the present invention not only suppresses yield point elongation due to normal temperature aging, but also suppresses a decrease in the total elongation value. This is presumed to be because the production of iron carbonitride at the grain boundary is suppressed, and is caused by the trapping of N and C at the Cu particle / Fe interface in the steel. Specifically, the deterioration amount of the total elongation is suppressed to 2% or less.
[0043]
Next, the reason for limiting the manufacturing method will be described.
The slab used for hot rolling is not particularly limited. That is, what was manufactured with the continuous casting slab, the thin slab caster, etc. should just be used. It is also compatible with processes such as continuous casting-direct rolling (CC-DR) in which hot rolling is performed immediately after casting.
[0044]
As for the hot rolling conditions, it is necessary to set the finishing temperature to (Ar3-100) ° C. or higher from the viewpoint of finely depositing Cu particles in the matrix during winding or cooling. If it is less than (Ar3-100) ° C., it causes coarse precipitation of Cu particles due to precipitation on the dislocation of Cu particles, and also causes a problem of plate thickness accuracy. The Ar3 temperature or higher is a more preferable range. The upper limit of the finishing temperature is not particularly defined, and the effects of the present invention can be obtained. However, in order to secure the r value, it is preferably set to 1000 ° C. or lower.
[0045]
After hot rolling, in order to precipitate Cu in the Fe matrix, a winding process is performed between 500 ° C. and 670 ° C. (the present invention according to (5) above), or a residence time between 550 and 720 ° C. is 170 s. It is necessary to cool so that it may become above (this invention which concerns on said (6)). When the coiling temperature exceeds 670 ° C., the Cu particles become coarse, and a Cu particle / Fe interface area of 1 [μm ² / μm ³ ] or more cannot be obtained, and as a result, room temperature non-aging property cannot be obtained. . When the winding temperature is less than 500 ° C., Cu particles do not precipitate within the normal winding time. Moreover, when the residence time between 550-720 degreeC is less than 170 s, precipitation of Cu will not occur.
[0046]
After the above cooling or winding, the sample is retained at 100 to 300 ° C. for 20 seconds or longer. Although the detailed mechanism is not clear, it is considered that segregation of solute C at the Cu particle / Fe interface occurs during this step, and is an essential step for obtaining suppression of yield point elongation during holding at room temperature. A temperature range of 100 to 300 ° C. is the temperature range for causing the above atom movement most rapidly in the crystal grains. If the residence time within this temperature range is less than 20 s, both BH properties and room temperature slow aging can be achieved. Can not. Therefore, the residence time range was limited to 20 s or more. In addition, from the viewpoint of obtaining better room temperature slow aging, holding for 60 seconds or more is more preferable. The upper limit of the residence time is not particularly defined, but if it exceeds 60000 s, precipitation of coarse carbonitrides at the grain boundary occurs, high BH may not be obtained, and the elongation value is significantly reduced. Is preferably 60000 s or less.
[0047]
After hot rolling, pickling may be performed as necessary, and then in-line or off-line temper rolling or leveling with a reduction rate of 3% or less, or cold rolling to a reduction rate of about 40% may be performed (the above ( The present invention according to 11)).
[0048]
Next, manufacturing conditions when a cold-rolled plate or a plated plate is used as the final product will be described.
The manufacturing conditions for the hot-rolled sheet as a material need not be specified, and may be performed according to a conventional method. Subsequently, a generally known treatment such as pickling is performed and cold rolling is performed. With regard to the cold rolling conditions, the number of rolling passes and the rolling reduction need not be specifically defined, and may be in accordance with ordinary methods. However, if the rolling reduction of cold rolling exceeds 90%, the load on the equipment becomes excessive, and the anisotropy of the mechanical properties of the product increases, so 90% or less is preferable.
[0049]
What is necessary is just to follow a conventional method about the heating rate in a continuous annealing process or a continuous annealing and a plating process. With respect to the maximum temperature achieved during the primary heat treatment, recrystallization may not be completed if the temperature is lower than the Ac1 temperature, and workability may be deteriorated, or cementite may not be sufficiently dissolved and BH may not be obtained. On the other hand, when the primary heat treatment temperature exceeds (Ac1 +200) ° C., the plate breaks in the plate and the productivity is greatly reduced. Therefore, the range of the maximum temperature reached in the primary heat treatment was limited to Ac1 to (Ac1 + 200) ° C.
[0050]
After the primary heat treatment, in order to precipitate Cu in the Fe matrix, it is necessary to perform a precipitation treatment in which the residence time between 550 and 720 ° C. is 170 s or more. When the residence time between 550 and 720 ° C. is less than 170 s, Cu does not precipitate, and it is difficult to ensure room temperature non-aging (the present invention according to (7) above).
[0051]
When hot dip galvanization is not performed, after the primary heat treatment and after cooling to 500 ° C., it may be cooled to any temperature from 500 ° C. to room temperature as the overaging treatment ((8) above) The present invention).
[0052]
When hot dip galvanizing is performed, after the primary heat treatment, a precipitation treatment with a residence time of at least 550 to 720 ° C. of 170 s or more is performed, galvanization is performed, and then a plating phase is alloyed as necessary. Although the conditions for galvanizing and alloying are not particularly defined, the immersion time in the plating bath and the holding time in the alloying furnace are each 40 s or less, more preferably from the viewpoint of suppressing the precipitation of added C at the grain boundaries. It is preferably 20 s or less (the present invention according to (9) and (10) above).
[0053]
Cooling is performed in which the residence time between 100 and 300 ° C. is 20 s or more after the primary heat treatment, after the heat treatment using the overaging zone, or after the plating treatment (including after the alloying treatment). In addition, when performing an overaging process, in order to ensure residence time in the above-mentioned temperature range, it is preferable on heat efficiency to set it as 100-300 degreeC and 20 s or more. In addition, the cooling process whose residence time between 100-300 degreeC is 20 s or more is an indispensable process in order to obtain yield-point elongation expression suppression during normal temperature holding | maintenance as mentioned above. The temperature range from 100 to 300 ° C. is the temperature range for causing the above atom movement most rapidly in the crystal grains. If the residence time within this temperature range is less than 20 s, both BH properties and room temperature non-aging properties can be achieved. Can not. Therefore, the residence time range was limited to 20 s or more. In addition, from the viewpoint of obtaining superior room temperature non-aging properties, holding for 60 seconds or more is more preferable. The upper limit of the residence time is not particularly defined, but if it exceeds 60000 s, precipitation of coarse carbonitrides at the grain boundary occurs, high BH may not be obtained, and the elongation value is significantly reduced. Is preferably 60000 s or less.
[0054]
Even when a cold-rolled sheet or a plated sheet is used as a final product, temper rolling or leveler processing is preferably performed in order to improve normal temperature slow aging and shape correction, and in a range of a rolling reduction of 3% or less. If it exceeds 3%, the amount of BH tends to decrease, so this is the upper limit (the present invention according to (11) above).
[0055]
The cold-rolled steel sheet of the present invention produced without undergoing a plating step or a plating alloying step after the primary heat treatment is suitable as various plating raw materials. The plating layer may be formed by either an electroplating method or a hot dipping method. Examples of the main component of plating include aluminum, zinc, chromium, tin, and nickel.
[0056]
The interface area between the Cu particles / Fe is determined by observing an image with a transmission electron microscope, or when the Cu precipitation phase is extremely fine, the average Cu particle diameter and precipitation within the measurement region are measured by an atom probe field ion microscope. A method of measuring and determining the volume fraction of an object is preferred. A method of measuring the amount of N or C segregated at the Cu particle / Fe interface by an atom probe field ion microscope is simple and suitable. Regarding the size of the Cu particles, those having a diameter of 1 nm or more are regarded as precipitates.
[0057]
【Example】
Next, the present invention will be described in detail with reference to examples.
Steels having the components shown in Table 1 were melted, and the hot rolling process was performed under the conditions shown in Table 2. As shown in Table 2, the slab heating temperature during hot rolling was 1050 to 1250 ° C., and the final thickness was 4 mm. All temper rolling was performed at an elongation of 1.0%. The steel plate thus obtained was subjected to a tensile test, a BH test, and a structure observation.
[0058]
Moreover, about the steel of the component shown in Table 1, a slab is reheated to 1050-1250 degreeC, it hot-rolls to the final board thickness of 4 mm with the hot rolling completion temperature of 840-930 degreeC, and it winds at 550-650 degreeC, After pickling the hot-rolled steel sheet thus obtained, cold working at a cold rolling rate of 70 to 85% and degreasing treatment, followed by continuous heat treatment and continuous galvanizing process under the conditions shown in Table 3 Went. All the temper rolling ratios were 1.0%. The steel plate thus obtained was subjected to a tensile test, a BH test, and a structure observation. Conditions for each test and observation are shown below.
[0059]
The tensile test was performed using a JIS No. 5 test piece under the condition of a strain rate of 10 ⁻³ / s. Material change during normal temperature holding was evaluated by comparing the tensile test results before and after accelerated aging at 100 ° C. × 1 hr. On the other hand, the tensile test for observing the amount of BH was performed using a JIS13B test piece under the condition of a strain rate of 10 ⁻³ / s. The pre-deformation amount in the BH test was 2%, the aging conditions corresponding to the paint baking treatment were 170 ° C. × 20 minutes, and the BH amount evaluated at the upper yield point at the time of re-tensioning was taken. The average crystal grain size of ferrite was measured according to the test method of JISG0552. The test results are shown in Table 4.
For burring workability (stretch flangeability), Japan Iron and Steel Federation Standard JFS T
Evaluation was performed by the hole expansion test method described in 1001-1996.
[0060]
Sample No. 1, No. 1 No. 8 is an example in which a sufficient amount of Cu particle / Fe interface area was not obtained when the amount of Cu was outside the proper range, and as a result, room temperature non-aging characteristics could not be obtained. Sample No. 4, no. No. 12 has a hot rolling coiling temperature or a residence time between 720 ° C. and 550 ° C. outside the proper range, so a sufficient amount of Cu particle / Fe interface area cannot be obtained, and as a result, room temperature non-aging characteristics can be obtained. This is an example that did not exist. Sample No. 5, no. No. 10 is an example in which a sufficient amount of C segregation at the Cu particle / Fe interface could not be obtained because the residence time between 300 and 100 ° C. was outside the proper range, and as a result, room temperature non-aging characteristics could not be obtained. is there. Sample No. 7, no. No. 14 is an example in which the amount of change ΔYP-EL in yield point elongation did not fall within 0.6% because the amount of N was excessive and the amount of BH was excessive with respect to the amount of Al.
[0061]
[Table 1]

[0062]
[Table 2]

[0063]
[Table 3]

[0064]
[Table 4]

[0065]
【The invention's effect】
The present invention has excellent molding limit values suitable for structural members such as automotive structural members, suspension members, panel members, electrical product inner and outer panel panels, buildings, etc. subjected to electrodeposition coating baking treatment, Strain age-hardening-type steel with little material deterioration during holding, excellent burring workability, and high strain-hardening ability can be provided at low cost, and is industrially valuable.

Claims

% By mass
C: 0.002 to 0.008%,
Si: 0.5% or less,
Mn: 0.1 to 3.0%
P: 0.1% or less,
S: 0.1% or less,
Cu: 0.6-3.0%
Ni: 0.1 to 3.0%,
Al: 0.01% or more,
N: 0.01% or less, N ^* defined by the following formula is 0 or less, the balance is Fe and inevitable impurities, and the area of the interface between Cu precipitates and Fe in the steel is per unit volume. 1 [μm ² / μm ³ ] or more, with ferrite as the maximum area ratio phase, and a change in yield point elongation ΔYP-El due to aging at 100 ° C. for 1 hour is 0.6% or less. Strain age hardened steel with excellent room temperature non-aging and burring workability.
N ^* = N-14 / 27 × Al

The strain age hardening type steel material excellent in normal temperature non-aging property and burring workability according to claim 1, further comprising one group or two or more groups of the following groups a to d in addition to the composition.
Group a: 0.005 to 1.0% of the total of one or more of Cr, Mo, and W.
Group b: 0.001 to 0.2% of the total of one or more of Nb, Ti, V, and Ta.
c group: B is 0.0003 to 0.010%.
d group: One or more of Ca, Mg, Zr, Ce, and REM in total 0.001 to 0.01%.

According to claim 1 or 2 C of segregation _{n excess} of the interface of Cu precipitates and Fe is equal to or is Cu precipitates and 0.05 per unit area of the interface of Fe [atoms / nm ^two] or more Strain age hardened steel with excellent non-aging at room temperature and burring workability.

A strain age hardening type steel material excellent in normal temperature non-aging property and burring workability, wherein the steel material according to any one of claims 1 to 3 is electroplated or hot dipped.

The slab comprising the chemical component according to claim 1 or 2 is heated, hot-rolled at a rolling end temperature (Ar3-100) ° C or higher, wound at 500 to 670 ° C, and then between 100 to 300 ° C. The strain aging hardening type steel material excellent in room temperature non-aging property and burring workability according to any one of claims 1 to 3, wherein the aging treatment or cooling in which the residence time is 20 s or more is performed. Method.

The slab comprising the chemical component according to claim 1 or 2 is heated, subjected to hot rolling at a rolling end temperature (Ar3-100) ° C. or higher, and cooling at a residence time of 550 to 720 ° C. of 170 s or longer. The aging treatment or cooling in which the residence time between 100 to 300 ° C is 20 s or more is then performed, and the room temperature non-aging property and burring workability are excellent in any one of claims 1 to 3 A method for producing a strain age-hardening steel.

After cold-rolling the steel material comprising the chemical component according to claim 1 or 2, the cold-rolled sheet is subjected to a primary heat treatment at a maximum temperature between Ac1 and (Ac1 +200) ° C, and then between 550 and 720 ° C. The aging treatment with a residence time of 170 s or more is performed, and the aging treatment or cooling with a residence time between 100 to 300 ° C of 20 s or more is further performed. A method for producing strain age-hardening type steels with excellent non-aging at room temperature and burring workability.

The non-aging property at room temperature according to claim 7, wherein after the primary heat treatment and the aging precipitation treatment, an overaging treatment is performed, and then an aging treatment or cooling in which a residence time between 100 to 300 ° C. is 20 s or more is performed. A method for producing strain age-hardening steel with excellent burring workability.

After cold-rolling the steel material comprising the chemical component according to claim 1 or 2, the cold-rolled sheet is subjected to primary heat treatment at a maximum temperature between Ac1 and (Ac1 +200) ° C, and then between 550 and 720 ° C. An aging precipitation treatment with a residence time of 170 s or more is performed, then a hot-dip galvanized layer is formed on the surface of the steel material, and an aging treatment or cooling with a residence time between 100 to 300 ° C. of 20 s or more is performed. The method of manufacturing the strain age hardening-type steel material excellent in normal temperature non-aging property and burring workability of any one of Claims 1-3.

After forming the hot-dip galvanized layer, alloying treatment is performed, and then the aging treatment or cooling in which the residence time between 100 to 300 ° C is 20 s or more is performed. A method for producing strain age-hardening steel with excellent burring workability.

The steel material manufactured by the method according to any one of claims 5 to 10 is subjected to temper rolling or leveler processing with an elongation of 3% or less, and is excellent in normal temperature non-aging and burring workability A method for producing strain age-hardening steel.