JP3821042B2 - High-formability high-tensile steel sheet with excellent strength stability and method for producing and processing the same - Google Patents
High-formability high-tensile steel sheet with excellent strength stability and method for producing and processing the same Download PDFInfo
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Description
【0001】
【発明の属する技術分野】
本発明は、自動車等の輸送機に使用される部材に適した、引張強度が550MPa以上の強度安定性に優れた高成形性高張力鋼板ならびにその製造方法および加工方法に関する。
【0002】
【従来の技術】
近年、自動車を代表する輸送機分野において、燃費向上を目的に車体の軽量化が検討されている。この車体軽量化の検討の一つとして、使用鋼板の高強度化が推進されている。従来より、高強度鋼板として、C,Mn,Siの固溶強化と、Ti,Nbの析出強化を複合した製造コストの低い鋼板が用いられてきたが、Cを0.12〜0.15%含有するため、鋼中にセメンタイトが多く析出し延性が乏しく、自動車用部材のような難加工材ではプレス割れを起こしていた。
【0003】
このようなプレス割れを回避する高成形性高張力熱延鋼板として、特開平6−172924号公報にはSiで炭化物析出を抑制するとともに、Cr添加量制限で低温変態相生成を抑制し、組織をベイニティックフェライト単相とし、さらにNi,Moを固溶強化元素として用いたTi添加高伸びフランジ加工性高張力熱延鋼板が開示されている。しかしながら、この技術の根幹をなすラス間に炭化物析出をともなわないラス状組織であるベイニティックフェライト組織では、Ti添加量を炭窒化物形成限界以下に制限していることから、Tiと結合しないCをベイニティックフェライト中に過飽和に固溶させなければならない。このため、炭化物析出駆動力が高い状態であることから添加成分の微妙な増減や熱延条件に対して炭化物析出の感受性が著しく強く、通常起こりうる幅方向の温度変動に対しても炭化物が容易に析出するようになり、幅方向で部分的に加工性が急激に劣化するのが現状である。
【0004】
また、特開平7−11382号公報には、Cと結合するTi,Nb量をCに対して原子比で0.5以上添加し、固溶Ti,Nbで熱間圧延後のフェライト核生成を抑制することで組織をアシキュラーフェライトとし、さらにCr,Moの固溶強化で強度を調整した高伸びフランジ性熱延鋼板が開示されている。しかしながら、この技術におけるアシキュラーフェライト組織の熱延鋼板は843MPaの強度で伸びが15%であり、伸びが高いことが要求される張り出し成形に対しては延性が未だ十分ではなく、このような特性の熱延鋼板に対し実際に張り出し成形を行うと割れが生じてしまう。さらに、ただ単に極低炭素鋼にTi、Nbを添加し、固溶Ti、Nb量を十分確保してもアシキュラーフェライト組織は得られないように、Ti、Nbによる組織のアシキュラー化効果は極めて小さく、この技術では多少の製造条件の変動でアシキュラーフェライトが得られなくなってしまう。
【0005】
一方で、高加工性と高強度化をTi,Nb,V,Moの微細化効果で実現する方法が特開平11−152544号公報に開示されている。しかし、この技術では粒径を2μm以下にすることから、伸びの劣化は避けられずやはり張り出し成形で割れが生じてしまう。また、粒径があまりにも微細なため、粒成長性が極めて大きく、通常起こりうる幅方向の熱延条件の変動で2μmを超える粒が部分的に生じて混粒組織となり、加工性が急激に劣化するのが現状である。
【0006】
また、特開2000−328186号公報には、Ti、Nbのオーステナイト細粒化と再結晶抑制により平均結晶粒を2.0〜10μmとし、かつフェライト面積率を95%以上とすることで、伸びフランジ性の優れた超微細フェライト組織鋼板が開示されている。しかしながら、この技術は未再結晶オーステナイトからフェライトへ変態させることを主眼としていることから、圧延後、巻取までの間の温度変動でオーステナイトの再結晶回復率が大きく変化するため材質変動が大きい。また、細粒はコイル内で粒成長の駆動力が大きいため部分的な粒成長すなわち混粒が生じやすく強度は安定しない。このように、安定した品質の製品を得ることは不可能である。さらに、この方法では不可避的に転位密度が高くなり、伸びの低下は避けられない。加えて、高々685MPaの強度を得るのにTiを0.32%添加しなければならず、圧延荷重が増大し、同一強度の鋼板と比べて板の形状が劣悪なものとなる。
【0007】
さらに、特開平6−200351号公報には、ポリゴナルフェライトに対するパーライトや低温変態相の面積比が15%以下でポリゴナルフェライト中にTiCが分散した組織を有し、かつ、Moの固溶強化で強度調整を行った伸びフランジ性が優れた高強度熱延鋼板が開示されている。しかしながら、この鋼板では、TiCの析出温度が狭範囲なため、幅方向センターでは加工性が良好でも、エッジでは規定の強度を下回り、延性が著しく劣化するのが現状である。
【0008】
【発明が解決しようとする課題】
このように従来技術では、加工性に優れた高張力鋼板が提案されているものの、コイル内の引張特性の変動が大きかったり、通常起こりうる製造条件の変動で加工性が劣化してしまう等、工業生産に適さないのは明らかである。
【0009】
本発明はかかる事情に鑑みてなされたものであって、当然起こり得る製造条件の変動に起因するコイル内強度変動、特に長手方向の強度変動が最小限であり、工業的に実用可能な強度安定性に優れた高成形性高張力鋼板ならびにその製造方法および加工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、従来技術では解決されない熱延ランナウトテーブルからコイラにかけてのセンターとエッジの冷却履歴の違いによる幅方向の材質変化や圧延速度の変化で生じるランナウトテーブル上のストリップの冷却速度変化などが原因となるコイル内の長手方向の材質変動、特に強度変動を小さくするために鋭意研究を重ねた結果、上記従来技術に示すラス状組織でもベイニティックフェライトでも2μm以下の微細粒でもなく、フェライト単相組織を、Tiと、MoおよびWの1種以上とを含む微細析出物で強化した鋼において、Ti、Mo、Wと結合しないCであるEx.Cが0.015%以下で、かつ、Mnを0.2≦Mn1.7−30×Ex.Cとすることにより材質変動、特に強度変動が低減されることを見出した。また、このような組織とすることにより、従来不可欠であったSiの添加を極力低減することが可能となり、表面性状を悪化させるSiに起因する赤スケールの生成も抑制することができることも見出した。
【0011】
本発明はこのような知見に基づいてなされたものであって、以下の(1)〜(7)を提供する。
【0014】
(1)質量%で、C:0.03〜0.15%、Mn≧0.2%、Si≦0.3%、P≦0.02%、Al≦0.1%、S≦0.01%、N≦0.01%、Ti:0.05〜0.35%を含み、かつMo≦0.6%、W≦1.5%から選ばれる1種以上を含み、MoおよびWがそれぞれ単独で含まれる場合には、Mo≧0.1%、W≧0.2%であり、残部がFeおよび不可避不純物からなり、以下の(1)式で示すEx.Cが0.015%以下であり、かつ、Mn≦1.7−30×Ex.Cを満たし、体積%で95%以上がフェライト組織であり、該フェライト組織に、Tiと、MoおよびWのうち1種以上とを含む10nm未満の析出物が分散してなることを特徴とする引張強度が550MPa以上の強度安定性に優れた高成形性高張力鋼板。
Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)…(1)
(ただし、(1)式の元素記号は各元素の質量%を示す。)
【0015】
(2)質量%で、C:0.03〜0.15%、Mn≧0.2%、Si≦0.3%、P≦0.02%、Al≦0.1%、S≦0.01%、N≦0.01%、Ti:0.05〜0.35%を含み、かつMo≦0.6%、W≦1.5%から選ばれる1種以上を含み、MoおよびWがそれぞれ単独で含まれる場合には、Mo≧0.1%、W≧0.2%であり、さらに、Cr≦0.5%、V≦0.1%、Nb≦0.08%、B≦0.001%の1種または2種以上を含み、残部がFeおよび不可避不純物からなり、以下の(1)式で示すEx.Cが0.015%以下であり、かつ、Mn≦1.7−30×Ex.Cを満たし、体積%で95%以上がフェライト組織であり、該フェライト組織に、Tiと、MoおよびWのうち1種以上とを含む10nm未満の析出物が分散してなることを特徴とする引張強度が550MPa以上の強度安定性に優れた高成形性高張力鋼板。
Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)…(1)
(ただし、(1)式の元素記号は各元素の質量%を示す。)
【0016】
(3)上記(1)または(2)において、板厚が2.0mm未満であることを特徴とする引張強度が550MPa以上の強度安定性に優れた高成形性高張力鋼板。
【0017】
(4)上記(1)または(2)の成分組成を有する鋼をオーステナイト単相域の温度に加熱後、熱間圧延を行うにあたり、800℃以上で仕上圧延を完了し、575〜675℃で巻取ることを特徴とする引張強度が550MPa以上の強度安定性に優れた高成形性高張力鋼板の製造方法。
【0018】
(5)上記(1)〜(3)のいずれかの鋼板からなる部材を準備する第1の工程と、前記部材にプレス成形を施して所望の形状のプレス成形品に加工する第2の工程とを有する高成形性高張力鋼板の加工方法。
【0019】
(6)上記(5)において、プレス成形品は、自動車用部品、特に自動車用足廻り部材である高成形性高張力鋼板の加工方法。
【0020】
(7)上記(1)から(3)のいずれかに記載の鋼板により製造された自動車用部品。
【0021】
このような構成の本発明によれば、(1)フェライト組織が形成され、γ→α変態直後の多量の固溶Cと偏析するMnの相互作用によるセメンタイトまたはパーライト等の製造熱履歴で形態が変化する粗大Fe炭化物の析出がないか、もしくは最小限に抑制されること、および(2)MoまたはWまたはMo+Wの作用によりランナウトテーブル上のγ→α変態が遅延され、広い温度域で安定的に析出するTiと、MoおよびWのうち1種以上とを含む微細炭化物が巻取り時に進行するフェライト変態とともに析出するようになることにより、鋼中Cを消費することで固溶Cが低減し、さらにはランナウトテーブル上の温度変化やコイル内変動が生じても組織変動が抑えられ、材質均一性に優れた鋼板が得られる。また、実質的にフェライト組織にTiと、MoおよびWのうち1種以上とを含む微細な炭化物が分散析出するため、高成形性でかつ高強度が実現される。
【0022】
さらに、本発明においては従来技術では所望の特性を得るために一定量以上必要であったSiを極力低減することが可能となり、表面性状を劣化させるSiに起因する赤スケールの生成を抑制することができる。
【0023】
【発明の実施の形態】
以下、本発明について具体的に説明する。
本発明に係る鋼板は、質量%で、C:0.03〜0.15%、Mn≧0.2%、Si≦0.3%、P≦0.02%、Al≦0.1%、S≦0.01%、N≦0.01%、Ti:0.05〜0.35%を含み、かつMo≦0.6%、W≦1.5%から選ばれる1種以上を含み、MoおよびWがそれぞれ単独で含まれる場合には、Mo≧0.1%、W≧0.2%であり、残部がFeおよび不可避不純物からなり、以下の(1)式で示すEx.Cが0.015%以下であり、かつ、Mn≦1.7−30×Ex.Cを満たし、体積%で95%以上がフェライト組織であり、該フェライト組織に、Tiと、MoおよびWのうち1種以上とを含む10nm未満の析出物が分散してなるものである。
Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)…(1)
(ただし、(1)式の元素記号は各元素の質量%を示す。)
【0024】
体積%で95%以上がフェライト組織と実質的にフェライト単相組織としたのは、複合組織では2種以上の組織形成を制御しなければならず、材質均一性を実現するのが困難であるのに対し、フェライト単相では複数の組織を同時に制御する困難性を解消することができるからである。たとえば、Fe炭化物はストリップやコイルの熱履歴により形態が変化し、これが多量に含まれていると材質変動の原因となる。
【0025】
具体的には断面組織観察などにより体積%で95%以上がフェライトとなっていればよい。好ましくは98%以上である。また、微細析出物以外の粗大なFe炭化物は体積%で1%未満であれば本発明の効果を損なうことがない。
【0026】
巻取り時にフェライト変態させるには、通常ランナウト冷却時に起こるγ→α変態を巻取りまで遅延させる必要がある。そこで、本発明ではMoまたはWを添加するか、またはWとMoを複合添加し、フェライト変態を遅延させる。
【0027】
マトリックスが実質的にフェライトからなる本発明の鋼板では、微細析出物により強度を担保する。この際の鋼板の強度は550MPa級以上であれば本発明の効果は得られるが、780MPa以上でその効果が一層高まり、980MPa以上でさらに一層その効果が高まる。一般にMoはMo炭化物を形成し、WはW炭化物を形成し、析出強化に寄与するが、Mo炭化物、W炭化物の析出速度は遅いため、MoまたはW単独では550MPa以上の高強度が実現しにくいうえに、巻取り後のコイル冷却時に析出するため、冷却速度の速いコイル外周部と中央部とでは強度が変化してしまう。そこで、巻取り前後までMoまたは/およびWを含む炭化物の析出を促進するために、Tiを添加し、析出物の析出速度を適切な値になるよう制御する。
【0028】
フェライト単相組織のマトリックスに微細析出物を分散させたのは、転位密度を上げて高強度を実現すると全伸びが低下するが、微細析出部での高強度化はこのような不都合が生じにくいからである。また、析出物の粒径を10nm未満としたのは、10nm以上になると550MPa以上の強度を得難くなるからである。また、10nm以上の析出物で強化しようとすると、析出物の体積率を多くしなければならず、必然的に析出物形成元素の添加量を上げなければならず、実際には析出物形成元素の添加量を上げたことにともなうフェライトの細粒化で強度を維持することになり、全伸びが低下してしまう。以上のことから本発明では微細析出物の粒径を10nm未満としたが、望ましくは5nm以下であり、さらに高強度が必要な場合には3nm以下とすることが望ましい。
【0029】
次に、上記成分組成について説明する。
C:0.03〜0.15%
Cは、Tiと、MoおよびWのうち1種以上とを含む炭化物として固定され、鋼の強度を担うのに必要不可欠な元素である。しかし、その含有量が0.03%未満では550MPa級以上の十分な強度は得られず、一方、0.15%を超えると粗大なFe炭化物が生成して延性が劣化する。そのため、C含有量を0.03〜0.15%とした。780MPa以上の強度を得るには0.035%以上であることが好ましい。
【0030】
N≦0.01%
Nは鋼中の不純物である。その含有量が0.01%を超えると延性を低下させる粗大な窒化物形成の原因となることから、0.01%以下とする。0.006%を超えるとTiを多量に添加しなければならなくなることから、0.006%以下が好ましい。
【0031】
Ti:0.05〜0.35%
TiはMoまたは/およびWとともに炭化物を形成し、鋼の強度を担う。しかし、0.05%未満では微細析出物量が少なくなり、高強度を実現することができなくなる。一方、0.35%を超えると伸び(EL)を低下させるベイニティックフェライトを生成しやすくなる。したがって、Ti含有量を0.05〜0.35%とした。
【0032】
Mo≦0.6%
Moは、上述したように、ランナウトテーブル上でのフェライト変態を抑制し、フェライト組織形成に対するランナウトテーブル上の熱履歴の影響を低減するともに、パーライト変態を抑制し、Tiとともに微細な炭化物を形成して鋼の高強度化に寄与する。しかし、0.6%以上添加すると低温変態相が多量に生じ、材質の安定性が劣化する。したがって、Mo含有量を0.6%以下とした。Wがともに含まれる場合には、Mo含有量の下限は存在しないが、Wが含まれない場合、十分な量のTiとMoとを含む炭化物を形成するために、Mo含有量を0.1%以上とする。
【0033】
W≦1.5%
Wは、上述したように、ランナウトテーブル上でのフェライト変態を抑制し、フェライト組織形成に対するランナウトテーブル上の熱履歴の影響を低減する。また、Tiとともに微細な炭化物を形成して鋼の高強度化に寄与する。しかし、Wが1.5%を超えると低温変態相が生成しやすくなり、材質の安定性が劣化する。したがって、W含有量を1.5%以下とした。Moがともに含まれる場合には、Wの下限は存在しないが、Moが含まれていない場合、Wが0.2%未満ではFe炭化物の析出を抑制することができなくなることから、W含有量を0.2%以上とする。
【0034】
C、Ti、Mo、Wは、0.5≦(C/12)/{(Ti/48)+(W/184)+(Mo/96)}≦1.5を満たすことが好ましい。これは、鋼中のCと(Ti+Mo+W)との原子数比(Wを含まない場合にはCと(Ti+Mo)との原子数比、Moを含まない場合にはCと(Ti+W)との原子数比)が0.5〜1.5であることを示しており、このような範囲となるように、C、Ti、Moまたは/およびWの含有量を調整することにより、Tiと、MoおよびWのうち1種以上とを含む炭化物が微細に析出しやすくなる。
【0035】
本発明において、例えば、TiとMoとで析出物を形成する場合、780MPa級の強度を得る場合、C:0.035〜0.055%、Ti:0.07〜0.09%、Mo:0.15〜0.25%であり、980MPa級の強度を得る場合、C:0.070〜0.080%、Ti:0.16〜0.2%、Mo:0.30〜0.40%が好ましい。
【0036】
Ex.C≦0.015%
Ex.CはTi、Mo、Wと結合しないCであり、Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)(ただし、式中の元素記号は各元素の質量%を示す。)で表される。このようなTi、Mo、Wと結合しないEx.Cが多量に存在すると、パーライトが生成する。つまり、Mo、W、およびNやSと結合していないTiはCと結合して析出物の析出とともにCを消費するが、Ti、Mo、Wと結合しないEx.Cは未変態γに濃化してパーライトを形成してしまう。したがって、Ex.Cをパーライトが生成しない0.015%以下とした。好ましくは0.010%以下、さらに好ましくは0.005%以下である。
【0037】
Mn:0.2〜1.7−30×Ex.C
Mnは固溶強化元素であり、十分な強度を確保するために、その含有量を0.2%以上とするが、一方において、本発明では、従来、固溶強化元素として積極的に用いられてきたMn量を制限することで強度の変動を低減することに成功した。すなわち、Mnは凝固時に偏析するため、Mnの濃化部にCが集まりやすく、たとえEx.Cを0.015%以下にしても、Mn量が多いとMn濃化部にパーライトが生成してしまい、巻取温度の変動に対する強度安定性が保てない。このため、本発明では、Mn含有量をEx.C量と関連付けて1.7−30×Ex.C以下とした。なお、Ex.C量の計算値が負値となる場合には、Mn量は1.7%以下とする。
【0038】
このことを実験結果で説明する。
C:0.03〜0.1%、Mn:0.25〜1.8%とし、Ti、Mo量を変化させた鋼を1250℃に加熱後、仕上げ温度910℃、巻き取り温度650℃と600℃の条件で熱間圧延を行い、1.6mmtの980MPa級熱延板を製造した。得られた熱延板の引張試験を行い、強度を評価した。このとき、650℃巻取の強度と600℃巻取の強度との強度差を計算した。これは、巻き取り温度が600℃から650℃の間で変動したときの材質変動を示すものである。この結果を図1に示す。図1において、650℃巻取と600℃巻取の強度差が30MPa以下のものを○、30MPaを超えるものを×とした。Mn含有量が上記範囲では、650℃巻取と600℃巻取での強度差が小さいことがわかる。
【0039】
次に、上記C、N、Ti、Mo、W、Mn以外のSi、P、Al、Sについて説明する。また、本発明の鋼板は、さらに、Cr≦0.5%、V≦0.1%、Nb≦0.08%を含んでもよく、これらについても以下に説明する。
【0025】
【0040】
Si≦0.3%
Siは固溶強化元素としてよく用いられてきた。しかしながら、Siは赤スケールを生成し、表面性状を劣化させるとともに、γ→α変態時に未変態γへのCの濃縮を促進し低温変態相を生成しやすくしてしまう。したがって、Si量は0.3%以下が好ましい。さらには、0.2%以下がより望ましく、より好ましくは0.1%以下である。
【0041】
P≦0.02%
Pは固溶強化元素であるが、0.02%を超えて添加されると粒界への著しい偏析を招き延性が劣化するので、0.02%以下が好ましい。
【0042】
S≦0.01%
Sは、TiSとして固定される。このためSは材質特性に有効に作用するTi量を低減させ、また延性も低下させることから、0.01%以下が好ましい。さらに好ましくは0.005%以下であり、一層好ましくは0.003%以下である。
【0043】
Al≦0.1%
鋼中Alは脱酸材として使用される。しかし、その含有量が0.1%を超えると鋼の延性低下を招くことから、0.1%以下が好ましい。さらに好ましくは0.06%以下である。
【0044】
Cr≦0.5%
CrはMo、Wとともに添加されると高温でのフェライト変態抑制効果が顕著となる。圧延後フェライト変態がランナウトテーブル上で著しく進行した場合、ストリップの幅方向の温度変動がそのままフェライト変態(ランナウトテーブル上の位置)タイミングのずれを招き、変態後のフェライト組織に影響を与える。すなわち、幅方向の中央部と端部とでは機械的性質が大きく異なることになる。Crはこのような幅方向の機械的性質の変動を抑制する効果を促進する。しかしながら、Crが0.5%を超えるとMnと同様に低延性の低温変態相が生成しやすくなる。したがって、Cr含有量は0.5%以下が好ましい。なお、このような効果をより顕著とするには0.04%以上が好ましい。
【0045】
V≦0.1%
Vは、炭化物、窒化物を形成し、鋼を強化する。しかし、その量が0.1%を超えると粗大な炭化物、窒化物ができて延性が劣化することから0.1%以下が好ましい。
【0046】
Nb≦0.08%
Nbは鋼を適度に微細化し、結晶粒形状を整粒化する目的で添加する。しかし、0.08%を超えると結晶粒の極度の微細化をもたらし、均一伸びが低下する傾向があることから0.08%以下が好ましい。
【0047】
なお、上記成分以外に、B≦0.001%、Ni≦0.5%、Cu≦0.5%、Ca≦0.005%、Sb≦0.005%、Co:0.01%を、耐食性、粒界酸化防止、縦割れ防止などを目的として添加してもよい。
【0048】
本発明において、板厚は2.0mm未満にすることが好ましい。本発明の効果は板厚が2.0mm未満の場合に顕著となるからである。これは、2mm未満の鋼板では、ストリップの進行性から、特に圧延前端水冷に制限があること、さらには加速圧延しなければならないことに起因して、コイル前後端で強度が異なりやすいからである。
【0049】
次に、以上のような本発明の鋼板の好ましい製造条件について述べる。
ここでは、上記成分組成を有する鋼をオーステナイト単相域の温度に加熱後、熱間圧延するにあたり、800℃以上で仕上圧延を完了し、575〜675℃で巻取る。
【0050】
仕上圧延温度:800℃以上
仕上圧延温度は材質均一化のために重要である。800℃未満では幅方向の温度変化で加工γの再結晶率が変化してしまい、変態組織に変動が認められるようになることから、仕上圧延温度を800℃以上とした。
【0051】
巻取り温度:575〜675℃
本発明鋼ではTiと、MoおよびWのうち1種以上とを含む炭化物の析出で強度を確保する巻取り温度を、上記炭化物の析出しやすい575〜675℃とした。また、本発明鋼では、Mo、Wによりフェライト変態が抑制され、Ex.CとMnが制御されていることから、コイル内の巻取温度変動に関わらず幅方向で組織は均一となり、フェライト変態直後にTiと、MoおよびWのうち1種以上とを含む炭化物が析出する。このため、強度、延性ともに安定する。
【0052】
本発明の高張力熱延鋼板には、表面に溶融亜鉛系めっき皮膜を形成し、溶融亜鉛系めっき鋼板としたものも含む。本発明の高張力熱延鋼板は良好な加工性を有することから、溶融亜鉛系めっき皮膜を形成しても良好な加工性を維持することができる。ここで、溶融亜鉛系めっきとは、亜鉛および亜鉛を主体とした溶融めっきであり、亜鉛の他にAl、Cr等の合金元素を含んだものを含む。このような溶融亜鉛系めっきを施した本発明の高張力熱延鋼板は、めっきままでもめっき後合金化処理を行ってもかまわない。めっき前焼鈍温度については、450℃未満ではめっきがつかず、750℃超えでは強度低下が生じやすい。そのため、焼鈍温度は450℃以上、750℃以下が好ましい。
【0053】
なお、本発明の熱延鋼板は、黒皮ままでも酸洗材でもその特性に差違はない。調質圧延についても通常行われているものであれば特に規定はない。また、上記溶融亜鉛めっきは酸洗後でも黒皮ままでも問題はない。亜鉛めっきについては電気めっきも可能である。化成処理についても特に問題はない。鋳造後直ちにもしくは補熱を目的とした加熱を施した後にそのまま熱間圧延を行う直送圧延を行っても本発明の効果に影響はない。さらに、粗圧延後に仕上圧延前で、圧延材を加熱しても、粗圧延後、圧延材を接合して行う連続圧延を行っても、さらには圧延材の加熱と連続圧延を同時に行っても本発明の効果は損なわれない。
【0054】
本発明の熱延鋼板は、表面性状と延性に優れ、コイル内材質変動も少ないのでこれをプレス成形した場合、その特質が活かされ、自動車用部材、特にサスペンションアーム等の足廻り部材のようなプレス時の断面形状が複雑な部材を良好な品質で製造することができ、特に、プレス成形品の軽量化に資することができる。以下に具体的に、本発明に係る熱延鋼板の加工方法、換言すればプレス成形品の製造方法について説明する。
【0055】
図2は、本発明に係る熱延鋼板の加工方法の作業フローの一例を示すフローチャートである。この作業フローは、通常、本発明に係る鋼板を製造することまたはその製造された鋼板を例えばコイルにして目的場所に搬送することを前工程としており、まず、本発明に係る熱延鋼板を準備することから始まる(S0、S1)。この鋼板に対してプレス加工を施す前に、鋼板に対して前処理的な加工を施すこともあれば(S2)、裁断機により所定の寸法や形状に加工することもある(S3)。前者のS2の工程では、例えば鋼板の幅方向の所定箇所に切り込みや穿孔を行い、引き続くプレス加工を終えた段階またはそのプレス加工の過程で、所定の寸法および形状のプレス成形品または被プレス加工部材として切り離すことができるようにしておく。後者のS3の工程では、最終的なプレス成形品の寸法、形状等を予め考慮して、所定の寸法および形状の鋼板部材に加工(したがって裁断)するようにしておく。その後、S2およびS3の工程を経由した部材には、プレス加工が施され、最終的に目的とする寸法・形状の所望のプレス成形品が製造される(S4)。このプレス加工は、通常は多段階で行われ、3段階以上7段階以下であることが多い。
【0056】
S4の工程は、S2およびS3の工程を経由した部材に対してさらに所定の寸法や形状に裁断する工程を含む場合もある。この場合の「裁断」という作業は、例えば、少なくともプレス加工の過程で、S2およびS3の工程を経由した部材の端部のような最終的なプレス成形品には不要部分を切り離す作業であっても構わないし、また、S2の工程で設けられた鋼板の幅方向の切り込みや穿孔に沿って被プレス加工部材を切り離す作業であっても構わない。
【0057】
なお、図2中、N1ないしN3は、鋼板、部材、プレス成形品を、機械的にあるいは作業員による搬送作業である場合がある。
【0058】
こうして製造されるプレス成形品は、必要に応じて次工程に送られる。次工程としては、例えば、プレス成形品にさらに機械加工を施し、寸法や形状を調整する工程、プレス成形品を所定場所に搬送し、格納する工程、プレス成形品に表面処理を施す工程、プレス成形品を用いて自動車のような目的物を組み立てる組立工程がある。
【0059】
図3は、図2に示した作業を実際に行う装置と鋼板、部材、プレス成形品の流れとの関係を示すブロック図である。この図においては、本発明に係る熱延鋼板はコイル状で準備されており、プレス加工機によりプレス成形品が製造される。プレス加工機は多段プレスを行う機種のものであるが、本件発明はこれに限定されない。
【0060】
プレス加工機の前段に、裁断機その他の前処理機械を設置する場合(図2の(a))もあれば、設置しない場合(図2の(b))もある。裁断機が設置される場合には、コイルから供給される長尺の本発明に係る鋼板から、必要な寸法又は形状の部材を裁断し、この部材がプレス加工機においてプレス加工され、所定のプレス成形品となる。鋼板の幅方向に切り欠きや穿孔を施す前処理機械が設置される場合には、プレス加工機においてその切り欠きや穿孔に沿って裁断が行われても構わない。前処理機械を設置しない場合には、プレス加工機において鋼板がプレス加工される過程で、裁断が行われ、最終的に所定の寸法、形状を有するプレス成形品が製造される。なお、図3における「裁断」の意味は、図2における裁断と同じである。
【0061】
こうして製造されるプレス成形品は、その原材料として表面性状と延性に優れ、コイル内材質変動も少ない本発明に係る鋼板を使用しているので、良好で均一な品質を有するに至り、かかるプレス成型品の製造歩留も高い。このような特長は、プレス成形品が自動車用部材、特にサスペンションアーム等の足廻り部材である場合に特に有用である。
【0062】
【実施例】
表1に示す化学成分を有する鋼を溶製し、加熱温度1250℃、仕上圧延温度約890℃、巻取温度約620℃で熱間圧延を行い、板厚が1.4mmおよび1.2mmの鋼板を作製した。得られた鋼板から作製した薄膜を透過型電子顕微鏡(TEM)によって析出物を観察した。Ti、W、Mo等の組成をTEMに装備されたエネルギー分散型X線分光装置(EDX)による分析から把握した。また、マトリックスの組織観察を走査型電子顕微鏡(SEM)により行った。
【0063】
さらに、得られた鋼板の長手方向中央部の一の幅方向中央部と幅方向端部の2ヶ所からJIS5号試験片を採取し引張特性およびその変化を調査した。引張特性として幅方向中央部の引張強度(TS)および伸び(El)を求め、引張特性の変化については、幅方向中央部と端部の材質差の絶対値で評価した。また、鋼板の伸びフランジ性(λ)を評価する穴広げ試験は、日本鉄鋼連盟規格に従って行った。これら特性結果を表2に示す。
【0064】
表1に示すように、No.1〜7は、Ex.C量、Mn量、その他の成分組成が本発明の範囲内であり、マトリックス組織がフェライトであり、Tiと、MoおよびWのうち1種以上とを含む10nm未満の析出物が分散してなる、本発明の範囲内の本発明例であり、表2に示すように、良好な強度安定性を得ることができた。
【0065】
一方、本発明の範囲を外れる比較例のNo.8、9では、Ex.C量およびMn量が本発明の範囲から外れ、本発明例のNo.1〜7と比較して強度が低く、引張特性のばらつきも大きかった。
【0066】
【表1】
【0067】
【表2】
【0068】
【発明の効果】
以上説明したように、本発明によれば、鋼の成分組成を適切に制御し、実質的にフェライト組織に特定組成のTiと、MoおよびWのうち1種以上とを含む炭化物を分散析出した構成としたので、γ→α変態直後の多量の固溶Cと偏析するMnの相互作用によるセメンタイトまたはパーライト等の製造熱履歴で形態が変化する粗大Fe炭化物の析出を最小限に抑えることができ、また、MoまたはWまたはMo+Wの作用によりランナウトテーブル上のγ→α変態が遅延され、広い温度域で安定的に析出するTiと、MoおよびWのうち1種以上とを含む微細炭化物が巻取り時に進行するフェライト変態とともに析出するようになるので、ランナウトテーブル上の温度変化やコイル内変動が生じても組織変動が抑えられ、材質均一性に優れた鋼板が得られる。また、実質的にフェライト組織にTiと、MoおよびWのうち1種以上とを含む微細な炭化物が分散析出するため、高成形性でかつ高強度が実現される。
【図面の簡単な説明】
【図1】Mn含有量とEx.C量の適正範囲を示す図。
【図2】本発明に係る熱延鋼板の加工方法の作業フローの一例を示すフローチャート。
【図3】図1に示した作業を実際に行う装置と鋼板、部材、プレス成形品の流れとの関係を示すブロック図。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a high-formability, high-tensile steel sheet suitable for members used in transportation equipment such as automobiles and having excellent strength stability with a tensile strength of 550 MPa or more, and a manufacturing method and a processing method thereof.
[0002]
[Prior art]
In recent years, in the field of transport aircraft, which is representative of automobiles, weight reduction of vehicle bodies has been studied for the purpose of improving fuel efficiency. As one of the examinations for reducing the weight of the vehicle body, increasing the strength of the steel sheet used is being promoted. Conventionally, as a high-strength steel plate, a steel plate having a low production cost in which solid solution strengthening of C, Mn, Si and precipitation strengthening of Ti, Nb has been used, but C is 0.12 to 0.15%. As a result, a large amount of cementite is precipitated in the steel, resulting in poor ductility, and press-cracking has occurred in difficult-to-process materials such as automobile members.
[0003]
As a high formability high tensile hot rolled steel sheet that avoids such press cracks, JP-A-6-172924 discloses that carbide precipitation is suppressed by Si, and low-temperature transformation phase generation is suppressed by limiting the amount of Cr added. Has disclosed a Ti-added high-stretch flangeability high-tensile hot-rolled steel sheet using Bainitic ferrite single phase and Ni and Mo as solid solution strengthening elements. However, in the bainitic ferrite structure, which is a lath-like structure that does not cause carbide precipitation between the laths that form the basis of this technology, the Ti addition amount is limited to the carbonitride formation limit or less, so it does not bond with Ti. C must be supersaturated in bainitic ferrite. For this reason, since the carbide precipitation driving force is high, the carbide precipitation is extremely sensitive to subtle increases and decreases in additive components and hot rolling conditions. In the present state, the workability is abruptly deteriorated partially in the width direction.
[0004]
In JP-A-7-11382, the amount of Ti and Nb bonded to C is added in an atomic ratio of 0.5 or more to C, and ferrite nucleation after hot rolling with solute Ti and Nb is performed. A high-stretch flanged hot-rolled steel sheet is disclosed in which the structure is changed to acicular ferrite and the strength is adjusted by solid solution strengthening of Cr and Mo. However, the hot rolled steel sheet having an acicular ferrite structure in this technique has a strength of 843 MPa and an elongation of 15%, and the ductility is not yet sufficient for stretch forming that requires high elongation. When the hot-rolled steel sheet is actually stretched and formed, cracking occurs. Furthermore, the addition of Ti and Nb to ultra-low carbon steel and the acicular effect of the structure by Ti and Nb is extremely high so that an acicular ferrite structure cannot be obtained even if sufficient amounts of solute Ti and Nb are secured. With this technique, acicular ferrite cannot be obtained with a slight change in manufacturing conditions.
[0005]
On the other hand, Japanese Patent Application Laid-Open No. 11-152544 discloses a method for realizing high workability and high strength by the effect of refinement of Ti, Nb, V, and Mo. However, since this technique makes the particle diameter 2 μm or less, deterioration of elongation is unavoidable, and cracks are also generated in the overhang molding. In addition, since the grain size is too fine, grain growth is extremely large, and grains that exceed 2 μm are partly formed due to fluctuations in the widthwise hot rolling conditions that usually occur, resulting in a mixed grain structure, and workability is drastically increased. The current situation is that it deteriorates.
[0006]
Japanese Patent Laid-Open No. 2000-328186 discloses that the average grain size is set to 2.0 to 10 μm and the ferrite area ratio is set to 95% or more by reducing the austenite of Ti and Nb and suppressing recrystallization. An ultrafine ferritic steel sheet having excellent flangeability is disclosed. However, since this technique mainly focuses on transformation from unrecrystallized austenite to ferrite, the material change is large because the recrystallization recovery rate of austenite greatly changes due to temperature fluctuations between rolling and winding. Further, since fine grains have a large driving force for grain growth in the coil, partial grain growth, that is, mixed grains are likely to occur, and the strength is not stable. Thus, it is impossible to obtain a product with stable quality. Furthermore, this method inevitably increases the dislocation density, and a reduction in elongation is unavoidable. In addition, to obtain a strength of at most 685 MPa, Ti must be added in an amount of 0.32%, the rolling load increases, and the shape of the plate becomes inferior compared to a steel plate of the same strength.
[0007]
Furthermore, Japanese Patent Laid-Open No. 6-200351 has a structure in which TiC is dispersed in polygonal ferrite with an area ratio of pearlite and low-temperature transformation phase to polygonal ferrite of 15% or less, and solid solution strengthening of Mo Discloses a high-strength hot-rolled steel sheet having excellent stretch flangeability after strength adjustment. However, in this steel sheet, since the precipitation temperature of TiC is in a narrow range, even if the workability is good at the center in the width direction, the edge is below the specified strength and the ductility is significantly deteriorated.
[0008]
[Problems to be solved by the invention]
In this way, in the conventional technology, although a high-tensile steel plate excellent in workability has been proposed, the variation in tensile properties in the coil is large, the workability deteriorates due to variations in production conditions that can normally occur, etc. Clearly it is not suitable for industrial production.
[0009]
The present invention has been made in view of such circumstances. Naturally, fluctuations in the strength in the coil caused by fluctuations in manufacturing conditions that can occur naturally, especially fluctuations in strength in the longitudinal direction, are minimal, and can be used for industrially stable strength. An object of the present invention is to provide a high-formability high-tensile steel sheet having excellent properties and a method for producing and processing the same.
[0010]
[Means for Solving the Problems]
The inventors of the present invention have not been solved by the prior art, such as a change in the cooling rate of the strip on the runout table caused by a change in material in the width direction or a change in rolling speed due to a difference in cooling history between the center and the edge from the hot rolled runout table to the coiler. As a result of intensive studies to reduce fluctuations in the longitudinal direction of the material in the coil, particularly strength fluctuations, the lath structure and bainitic ferrite shown in the above prior art are not fine grains of 2 μm or less, In a steel in which a ferrite single-phase structure is strengthened with fine precipitates containing Ti and one or more of Mo and W, Ex. C is 0.015% or less, and Mn is 0.2 ≦ Mn1.7-30 × Ex. It was found that the material variation, particularly strength variation, is reduced by setting C. In addition, it has also been found that such a structure makes it possible to reduce the addition of Si, which has been indispensable in the past, as much as possible, and to suppress the generation of red scale due to Si that deteriorates the surface properties. .
[0011]
The present invention has been made based on such findings, and the following (1) to (7)I will provide a.
[0014]
(1) By mass%, C: 0.03 to 0.15%, Mn ≧ 0.2%, Si ≦ 0.3%, P ≦ 0.02%, Al ≦ 0.1%, S ≦ 0. 01%, N ≦ 0.01%, Ti: 0.05 to 0.35%, and one or more selected from Mo ≦ 0.6% and W ≦ 1.5%, and Mo and W are When each of them is contained alone, Mo ≧ 0.1% and W ≧ 0.2%, the balance is made of Fe and inevitable impurities, and Ex. C is 0.015% or less, and Mn ≦ 1.7-30 × Ex. C is satisfied, and 95% or more by volume% is a ferrite structure, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed in the ferrite structure. A high-formability, high-tensile steel plate with excellent strength stability with a tensile strength of 550 MPa or more.
Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184) (1)
(However, the element symbol in the formula (1) indicates mass% of each element.)
[0015]
(2) By mass%, C: 0.03-0.15%, Mn ≧ 0.2%, Si ≦ 0.3%, P ≦ 0.02%, Al ≦ 0.1%, S ≦ 0. 01%, N ≦ 0.01%, Ti: 0.05 to 0.35%, and one or more selected from Mo ≦ 0.6% and W ≦ 1.5%, and Mo and W are When each is included alone, Mo ≧ 0.1%, W ≧ 0.2%, and Cr ≦ 0.5%, V ≦ 0.1%, Nb ≦ 0.08%, B ≦ Ex. 0.001% containing one or more, the balance consisting of Fe and unavoidable impurities, Ex. C is 0.015% or less, and Mn ≦ 1.7-30 × Ex. C is satisfied, and 95% or more by volume% is a ferrite structure, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed in the ferrite structure. A high-formability, high-tensile steel plate with excellent strength stability with a tensile strength of 550 MPa or more.
Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184) (1)
(However, the element symbol in the formula (1) indicates mass% of each element.)
[0016]
(3) Above (1)Or (2)A high-formability, high-tensile steel sheet excellent in strength stability with a tensile strength of 550 MPa or more, wherein the plate thickness is less than 2.0 mm.
[0017]
(4) Above (1)Or (2)In the hot rolling after heating the steel having the component composition to the temperature of the austenite single phase region, the finish rolling is completed at 800 ° C. or higher, and the tensile strength is 550 MPa. The manufacturing method of the high formability high tension steel plate excellent in the above strength stability.
[0018]
(5) Above (1)-(3A high-formability high-tensile steel plate having a first step of preparing a member made of any one of the steel plates, and a second step of pressing the member into a press-formed product having a desired shape. Processing method.
[0019]
(6)the above(5), The press-formed product is a method for processing a high-formability high-tensile steel sheet, which is an automobile part, particularly an automobile suspension member.
[0020]
(7) From (1) to (3) An automotive part manufactured from the steel sheet according to any of the above.
[0021]
According to the present invention having such a configuration, (1) a ferrite structure is formed, and the form is formed by a production heat history of cementite or pearlite due to the interaction between a large amount of solute C immediately after the γ → α transformation and segregated Mn. There is no or minimal precipitation of changing coarse Fe carbide, and (2) γ → α transformation on the run-out table is delayed by the action of Mo or W or Mo + W, and stable over a wide temperature range As the fine carbides containing Ti and one or more of Mo and W are precipitated together with the ferrite transformation that proceeds during winding, solid solution C is reduced by consuming C in the steel. Furthermore, even if a temperature change on the run-out table or a fluctuation in the coil occurs, the structural fluctuation is suppressed, and a steel sheet having excellent material uniformity can be obtained. In addition, since fine carbides containing Ti and one or more of Mo and W are substantially dispersed and precipitated in the ferrite structure, high formability and high strength are realized.
[0022]
Furthermore, in the present invention, it is possible to reduce Si as much as possible to obtain a desired characteristic in the prior art as much as possible, and suppress the generation of red scale due to Si that deteriorates the surface properties. Can do.
[0023]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be specifically described.
The steel sheet according to the present invention is in mass%, C: 0.03 to 0.15%, Mn ≧ 0.2%, Si ≦ 0.3%, P ≦ 0.02%, Al ≦ 0.1%, Including S ≦ 0.01%, N ≦ 0.01%, Ti: 0.05 to 0.35%, and including one or more selected from Mo ≦ 0.6% and W ≦ 1.5%, When Mo and W are contained alone, Mo ≧ 0.1% and W ≧ 0.2%, the balance is made of Fe and inevitable impurities, and Ex. C is 0.015% or less, and Mn ≦ 1.7-30 × Ex. C is satisfied, and 95% or more by volume% is a ferrite structure, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed in the ferrite structure.
Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184) (1)
(However, the element symbol in the formula (1) indicates mass% of each element.)
[0024]
More than 95% by volume is ferrite structure and ferrite single phase structure substantiallyThis is because, in the composite structure, two or more types of structures must be controlled and it is difficult to achieve material uniformity, whereas in the ferrite single phase, it is difficult to control multiple structures simultaneously. It is because it can be eliminated. For example, Fe carbides change in shape due to the thermal history of strips and coils, and if they are contained in a large amount, they cause material fluctuations.
[0025]
IngredientsPhysically, cross-sectional structure observation etc.RIt is only necessary that 95% or more by volume is ferrite. Preferably it is 98% or more. Further, if the coarse Fe carbide other than fine precipitates is less than 1% by volume, the effect of the present invention is not impaired.
[0026]
In order to transform the ferrite during winding, it is necessary to delay the γ → α transformation that normally occurs during run-out cooling until the winding. Therefore, in the present invention, Mo or W is added, or W and Mo are added in combination to delay the ferrite transformation.
[0027]
In the steel sheet of the present invention in which the matrix is substantially composed of ferrite, the strength is ensured by fine precipitates. In this case, if the strength of the steel sheet is 550 MPa or more, the effect of the present invention can be obtained. However, the effect is further enhanced at 780 MPa or more, and the effect is further enhanced at 980 MPa or more. In general, Mo forms Mo carbide and W forms W carbide, which contributes to precipitation strengthening. However, because the precipitation rate of Mo carbide and W carbide is slow, it is difficult to achieve high strength of 550 MPa or more with Mo or W alone. In addition, since it precipitates when the coil is cooled after winding, the strength changes between the outer peripheral portion and the central portion of the coil where the cooling rate is fast. Therefore, in order to promote precipitation of carbides containing Mo or / and W before and after winding, Ti is added and the precipitation rate of the precipitates is controlled to an appropriate value.
[0028]
The reason why fine precipitates are dispersed in the matrix of ferrite single phase structure is that when the dislocation density is increased and high strength is achieved, the total elongation is reduced. Because. Moreover, the reason why the particle size of the precipitate is less than 10 nm is that it becomes difficult to obtain a strength of 550 MPa or more when the particle diameter is 10 nm or more. In addition, when trying to strengthen with precipitates of 10 nm or more, the volume fraction of the precipitates must be increased, and the amount of the precipitate-forming elements must be increased. The strength is maintained by making the ferrite finer with increasing the amount of added, and the total elongation is reduced. From the above, in the present invention, the particle size of the fine precipitate is set to less than 10 nm, but is preferably 5 nm or less, and more preferably 3 nm or less when higher strength is required.
[0029]
Next, the component composition will be described.
C: 0.03-0.15%
C is fixed as a carbide containing Ti and one or more of Mo and W, and is an element indispensable for bearing the strength of steel. However, if the content is less than 0.03%, sufficient strength of 550 MPa class or more cannot be obtained. On the other hand, if the content exceeds 0.15%, coarse Fe carbide is generated and ductility deteriorates. Therefore, the C content is set to 0.03 to 0.15%. In order to obtain a strength of 780 MPa or more, it is preferably 0.035% or more.
[0030]
N ≦ 0.01%
N is an impurity in the steel. If the content exceeds 0.01%, coarse nitride formation that lowers the ductility is caused, so the content is made 0.01% or less. If it exceeds 0.006%, a large amount of Ti must be added, so 0.006% or less is preferable.
[0031]
Ti: 0.05 to 0.35%
Ti forms carbide with Mo or / and W, and bears the strength of steel. However, if it is less than 0.05%, the amount of fine precipitates decreases, and high strength cannot be realized. On the other hand, if it exceeds 0.35%, bainitic ferrite that lowers the elongation (EL) tends to be generated. Therefore, the Ti content is set to 0.05 to 0.35%.
[0032]
Mo ≦ 0.6%
As described above, Mo suppresses ferrite transformation on the runout table, reduces the influence of thermal history on the runout table on ferrite structure formation, suppresses pearlite transformation, and forms fine carbide with Ti. This contributes to increasing the strength of steel. However, when 0.6% or more is added, a large amount of low-temperature transformation phase is generated, and the stability of the material is deteriorated. Therefore, the Mo content is set to 0.6% or less. When W is included together, there is no lower limit of the Mo content. However, when W is not included, in order to form a carbide containing a sufficient amount of Ti and Mo, the Mo content is set to 0.1%. % Or more.
[0033]
W ≦ 1.5%
As described above, W suppresses the ferrite transformation on the runout table and reduces the influence of the thermal history on the runout table on the ferrite structure formation. In addition, fine carbides are formed together with Ti, contributing to high strength of steel. However, if W exceeds 1.5%, a low temperature transformation phase is likely to be generated, and the stability of the material is deteriorated. Therefore, the W content is set to 1.5% or less. When Mo is included together, there is no lower limit of W. However, when Mo is not included, it is impossible to suppress precipitation of Fe carbide if W is less than 0.2%. Is 0.2% or more.
[0034]
C, Ti, Mo, and W preferably satisfy 0.5 ≦ (C / 12) / {(Ti / 48) + (W / 184) + (Mo / 96)} ≦ 1.5. This is the atomic ratio between C and (Ti + Mo + W) in steel (the atomic ratio between C and (Ti + Mo) when W is not included, and the atomic ratio between C and (Ti + W) when Mo is not included. Number ratio) is 0.5 to 1.5, and by adjusting the content of C, Ti, Mo or / and W so as to be in such a range, Ti and Mo And carbides containing one or more of W easily precipitate finely.
[0035]
In the present invention, for example, when forming a precipitate with Ti and Mo, when obtaining a strength of 780 MPa class, C: 0.035 to 0.055%, Ti: 0.07 to 0.09%, Mo: 0.15 to 0.25%, and when obtaining a strength of 980 MPa class, C: 0.070 to 0.080%, Ti: 0.16 to 0.2%, Mo: 0.30 to 0.40 % Is preferred.
[0036]
Ex. C ≦ 0.015%
Ex. C is C which is not bonded to Ti, Mo, W, Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184)(However, the element symbols in the formula indicate the mass% of each element.)It is represented by Ex. That does not bond to such Ti, Mo, W. When C is present in a large amount, pearlite is generated. That is, Ti that is not bonded to Mo, W, and N or S binds to C and consumes C along with precipitation of precipitates, but does not bond to Ti, Mo, and W. Ex. C concentrates to untransformed γ to form pearlite. Therefore, Ex. C was set to 0.015% or less at which pearlite was not generated. Preferably it is 0.010% or less, More preferably, it is 0.005% or less.
[0037]
Mn: 0.2 to 1.7-30 × Ex. C
Mn is a solid solution strengthening element, and its content is 0.2% or more in order to ensure sufficient strength. On the other hand, in the present invention, Mn has been actively used as a solid solution strengthening element. By limiting the amount of Mn that has been produced, we succeeded in reducing fluctuations in strength. That is, since Mn segregates during solidification, C tends to collect in the Mn concentrated portion. Even if C is 0.015% or less, if the amount of Mn is large, pearlite is generated in the Mn-concentrated portion, and the strength stability against fluctuations in the winding temperature cannot be maintained. For this reason, in the present invention, the Mn content is set to Ex. In association with C amount, 1.7-30 × Ex. C or less. Ex. When the calculated value of the C amount is a negative value, the Mn amount is 1.7% or less.
[0038]
This will be explained by experimental results.
C: 0.03 to 0.1%, Mn: 0.25 to 1.8%, and after changing the amount of Ti and Mo to 1250 ° C, the finishing temperature is 910 ° C and the winding temperature is 650 ° C. Hot rolling was performed at 600 ° C. to produce a 1.6 mmt 980 MPa class hot rolled sheet. The obtained hot-rolled sheet was subjected to a tensile test to evaluate the strength. At this time, the strength difference between the strength of 650 ° C. winding and the strength of 600 ° C. winding was calculated. This shows the material fluctuation when the winding temperature fluctuates between 600 ° C. and 650 ° C. The result is shown in FIG. In FIG. 1, the case where the strength difference between the 650 ° C. winding and the 600 ° C. winding is 30 MPa or less is indicated by “◯”, and the case where the strength difference exceeds 30 MPa is indicated by “X”. It can be seen that when the Mn content is in the above range, the difference in strength between the 650 ° C. winding and the 600 ° C. winding is small.
[0039]
next,Si, P, Al, and S other than C, N, Ti, Mo, W, and Mn will be described. Further, the steel sheet of the present invention may further contain Cr ≦ 0.5%, V ≦ 0.1%, and Nb ≦ 0.08%, which will be described below.
[0025]
[0040]
Si ≦ 0.3%
Si has often been used as a solid solution strengthening element. However, Si generates a red scale, which deteriorates the surface properties and promotes the concentration of C into the untransformed γ during the γ → α transformation, thereby facilitating the formation of a low temperature transformation phase. Therefore, the Si content is preferably 0.3% or less. Furthermore, 0.2% or less is more desirable, More preferably, it is 0.1% or less.
[0041]
P ≦ 0.02%
P is a solid solution strengthening element, but if added over 0.02%, significant segregation to the grain boundary is caused and the ductility deteriorates, so 0.02% or less is preferable.
[0042]
S ≦ 0.01%
S is fixed as TiS. For this reason, S is preferably 0.01% or less because it reduces the amount of Ti that effectively acts on the material properties and also reduces the ductility. More preferably, it is 0.005% or less, More preferably, it is 0.003% or less.
[0043]
Al ≦ 0.1%
Al in steel is used as a deoxidizer. However, if its content exceeds 0.1%, the ductility of the steel is reduced, so 0.1% or less is preferable. More preferably, it is 0.06% or less.
[0044]
Cr ≦ 0.5%
When Cr is added together with Mo and W, the effect of suppressing ferrite transformation at a high temperature becomes remarkable. When the ferrite transformation after rolling proceeds remarkably on the runout table, the temperature variation in the width direction of the strip causes the deviation of the ferrite transformation (position on the runout table) as it is, and affects the ferrite structure after the transformation. That is, the mechanical properties greatly differ between the central portion and the end portion in the width direction. Cr promotes the effect of suppressing such fluctuations in the mechanical properties in the width direction. However, if Cr exceeds 0.5%, a low-ductility low-temperature transformation phase is likely to be generated as in the case of Mn. Therefore, the Cr content is preferably 0.5% or less. In addition, in order to make such an effect more remarkable, 0.04% or more is preferable.
[0045]
V ≦ 0.1%
V forms carbides and nitrides and strengthens the steel. However, if the amount exceeds 0.1%, coarse carbides and nitrides are formed and ductility deteriorates, so 0.1% or less is preferable.
[0046]
Nb ≦ 0.08%
Nb is added for the purpose of appropriately refining the steel and adjusting the crystal grain shape. However, if it exceeds 0.08%, the crystal grains are extremely refined and uniform elongation tends to decrease, so 0.08% or less is preferable.
[0047]
In addition,In addition to the above ingredients,B ≦ 0.001%, Ni ≦ 0.5%, Cu ≦ 0.5%, Ca ≦ 0.005%, Sb ≦ 0.005%, Co: 0.01%May be added for the purpose of corrosion resistance, prevention of grain boundary oxidation, prevention of vertical cracks, and the like.
[0048]
In the present invention, the plate thickness is preferably less than 2.0 mm. This is because the effect of the present invention becomes significant when the plate thickness is less than 2.0 mm. This is because in steel sheets of less than 2 mm, the strength at the front and rear ends of the coil is likely to be different due to the progress of the strip, especially due to the fact that there is a limit to water cooling at the front end of rolling, and further that accelerated rolling is required. .
[0049]
Next, preferable production conditions for the steel sheet of the present invention as described above will be described.
Here, when the steel having the above composition is heated to the temperature of the austenite single phase region and then hot-rolled, finish rolling is completed at 800 ° C. or higher and wound at 575 to 675 ° C.
[0050]
Finish rolling temperature: 800 ℃ or more
The finishing rolling temperature is important for making the material uniform. If the temperature is less than 800 ° C., the recrystallization rate of the processed γ changes due to the temperature change in the width direction, and fluctuations are observed in the transformation structure. Therefore, the finish rolling temperature is set to 800 ° C. or more.
[0051]
Winding temperature: 575-675 ° C
In the steel of the present invention, the coiling temperature at which strength is ensured by precipitation of carbide containing Ti and one or more of Mo and W was set to 575 to 675 ° C. at which the carbide is likely to precipitate. In the steel of the present invention, ferrite transformation is suppressed by Mo and W, and Ex. Since C and Mn are controlled, the structure becomes uniform in the width direction regardless of the coiling temperature variation in the coil, and carbides containing Ti and one or more of Mo and W are precipitated immediately after the ferrite transformation. To do. For this reason, both strength and ductility are stable.
[0052]
The high-tensile hot-rolled steel sheet of the present invention includes a hot-dip galvanized steel sheet that has a hot-dip galvanized film formed on the surface. Since the high-tensile hot-rolled steel sheet of the present invention has good workability, good workability can be maintained even when a hot-dip galvanized film is formed. Here, the hot dip galvanizing is hot dip plating mainly composed of zinc and zinc, and includes those containing alloy elements such as Al and Cr in addition to zinc. The high-tensile hot-rolled steel sheet of the present invention subjected to such hot-dip galvanizing may be subjected to an alloying treatment after plating or as it is plated. As for the pre-plating annealing temperature, if the temperature is lower than 450 ° C., the plating cannot be applied, and if it exceeds 750 ° C., the strength tends to decrease. Therefore, the annealing temperature is preferably 450 ° C. or higher and 750 ° C. or lower.
[0053]
In addition, the hot-rolled steel sheet of the present invention has no difference in its characteristics whether it is black or pickled. There is no particular restriction on temper rolling as long as it is usually performed. Moreover, the hot dip galvanization has no problem even after pickling or as it is black. For galvanization, electroplating is also possible. There is no particular problem with chemical conversion treatment. The effect of the present invention is not affected even if direct feed rolling, in which hot rolling is performed directly after casting or after heating for the purpose of supplementary heating, is performed. Furthermore, even if the rolled material is heated after the rough rolling and before the finish rolling, the continuous rolling performed by joining the rolled material after the rough rolling may be performed, or the heating and continuous rolling of the rolled material may be performed simultaneously. The effect of the present invention is not impaired.
[0054]
The hot-rolled steel sheet of the present invention is excellent in surface properties and ductility and has little fluctuation in the material in the coil, so when it is press-molded, its characteristics are utilized, such as automobile parts, particularly suspension members such as suspension arms. A member having a complicated cross-sectional shape at the time of pressing can be manufactured with good quality, and in particular, can contribute to weight reduction of a press-formed product. Hereinafter, a method for processing a hot-rolled steel sheet according to the present invention, in other words, a method for manufacturing a press-formed product will be described.
[0055]
FIG. 2 is a flowchart showing an example of a work flow of the method for processing a hot-rolled steel sheet according to the present invention. This work flow usually has a pre-process of manufacturing a steel plate according to the present invention or transporting the manufactured steel plate to a destination place as a coil, for example. First, a hot-rolled steel plate according to the present invention is prepared. (S0, S1). Before pressing the steel sheet, the steel sheet may be pre-processed (S2), or may be processed into a predetermined size or shape by a cutting machine (S3). In the former step S2, for example, cutting or drilling is performed at a predetermined position in the width direction of the steel sheet, and a press-formed product having a predetermined size and shape or pressed processing is performed at the stage where the subsequent press processing is completed or in the process of the press processing. It can be separated as a member. In the latter step of S3, the final press-molded product is processed (and thus cut) into a steel plate member having a predetermined size and shape in consideration of the size and shape of the final press-formed product in advance. Thereafter, the member that has undergone the steps S2 and S3 is subjected to press working, and finally a desired press-formed product having a desired size and shape is manufactured (S4). This press working is usually performed in multiple stages, and often has 3 stages or more and 7 stages or less.
[0056]
The step S4 may include a step of further cutting the member that has passed through the steps S2 and S3 into a predetermined size and shape. The operation of “cutting” in this case is, for example, an operation of cutting an unnecessary portion in a final press-formed product such as an end portion of a member that has passed through steps S2 and S3 at least in the process of pressing. Alternatively, it may be an operation of cutting the member to be pressed along the cutting or perforation in the width direction of the steel plate provided in the step S2.
[0057]
In FIG. 2, N1 to N3 may be a work of conveying a steel plate, a member, or a press-formed product mechanically or by an operator.
[0058]
The press-formed product manufactured in this way is sent to the next step as necessary. As the next process, for example, a further process is performed on the press-molded product to adjust dimensions and shape, a process of transporting and storing the press-molded product to a predetermined place, a process of subjecting the press-molded product to surface treatment, a press There is an assembly process for assembling an object such as an automobile using a molded product.
[0059]
FIG. 3 is a block diagram showing the relationship between the apparatus that actually performs the operation shown in FIG. 2 and the flow of the steel plate, member, and press-formed product. In this figure, the hot-rolled steel sheet according to the present invention is prepared in a coil shape, and a press-formed product is manufactured by a press machine. The press machine is of a type that performs multi-stage pressing, but the present invention is not limited to this.
[0060]
In some cases, a cutting machine or other pre-processing machine is installed in the front stage of the press machine (FIG. 2A), and in some cases, it is not installed (FIG. 2B). When a cutting machine is installed, a member having a required size or shape is cut from a long steel sheet according to the present invention supplied from a coil, and this member is pressed by a press machine, and a predetermined press It becomes a molded product. In the case where a pre-processing machine that performs notches and perforations in the width direction of the steel sheet is installed, the press machine may cut along the notches and perforations. When the pretreatment machine is not installed, cutting is performed in the process of pressing the steel plate in the press machine, and finally a press-formed product having a predetermined size and shape is manufactured. The meaning of “cutting” in FIG. 3 is the same as the cutting in FIG.
[0061]
The press-molded product manufactured in this way uses the steel sheet according to the present invention, which has excellent surface properties and ductility as the raw material, and has little fluctuation in the material in the coil, so that it has good and uniform quality. The production yield of goods is also high. Such a feature is particularly useful when the press-formed product is a member for an automobile, particularly a suspension member such as a suspension arm.
[0062]
【Example】
Steel having the chemical components shown in Table 1 is melted and hot-rolled at a heating temperature of 1250 ° C., a finish rolling temperature of about 890 ° C., and a winding temperature of about 620 ° C., and the plate thickness is 1.4 mm and 1.2 mm. A steel plate was produced. The thin film produced from the obtained steel plate was observed with a transmission electron microscope (TEM). The composition of Ti, W, Mo, etc. was grasped from analysis by an energy dispersive X-ray spectrometer (EDX) equipped in the TEM. Moreover, the structure | tissue observation of the matrix was performed with the scanning electron microscope (SEM).
[0063]
Furthermore, JIS No. 5 test specimens were collected from two places, one central part in the width direction and one end part in the width direction, of the central part in the longitudinal direction of the obtained steel sheet, and the tensile properties and changes thereof were investigated. The tensile strength (TS) and elongation (El) at the center in the width direction were obtained as tensile properties, and the change in tensile properties was evaluated by the absolute value of the material difference between the center and the end in the width direction. Moreover, the hole expansion test for evaluating the stretch flangeability (λ) of the steel sheet was performed according to the Japan Iron and Steel Federation standard. These characteristic results are shown in Table 2.
[0064]
As shown in Table 1, no. 1-7 are Ex. The amount of C, the amount of Mn, and other component compositions are within the scope of the present invention, the matrix structure is ferrite, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed. These are examples of the present invention within the scope of the present invention, and as shown in Table 2, good strength stability could be obtained.
[0065]
On the other hand, No. of a comparative example out of the scope of the present invention. 8 and 9, Ex. The amount of C and the amount of Mn deviate from the scope of the present invention. Compared with 1-7, the strength was low and the variation in tensile properties was also large.
[0066]
[Table 1]
[0067]
[Table 2]
[0068]
【The invention's effect】
As described above, according to the present invention, the steel component composition is appropriately controlled, and a carbide containing a specific composition of Ti and one or more of Mo and W is substantially dispersed and precipitated in the ferrite structure. Because of the structure, it is possible to minimize the precipitation of coarse Fe carbides whose morphology changes due to the production heat history of cementite or pearlite due to the interaction between a large amount of solute C immediately after the γ → α transformation and segregated Mn. Further, the γ → α transformation on the run-out table is delayed by the action of Mo or W or Mo + W, and Ti that precipitates stably in a wide temperature range, and fine carbide containing at least one of Mo and W are wound. Steel that has excellent material uniformity because it will precipitate together with the ferrite transformation that progresses at the time of machining, and even if temperature changes on the runout table or fluctuations in the coil occur, structural changes are suppressed. It is obtained. In addition, since fine carbides containing Ti and one or more of Mo and W are substantially dispersed and precipitated in the ferrite structure, high formability and high strength are realized.
[Brief description of the drawings]
FIG. 1 shows Mn content and Ex. The figure which shows the appropriate range of C amount.
FIG. 2 is a flowchart showing an example of a work flow of a method for processing a hot-rolled steel sheet according to the present invention.
FIG. 3 is a block diagram showing the relationship between the apparatus that actually performs the work shown in FIG. 1 and the flow of steel plates, members, and press-formed products.
Claims (7)
Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)…(1)
(ただし、(1)式の元素記号は各元素の質量%を示す。) In mass%, C: 0.03 to 0.15%, Mn ≧ 0.2%, Si ≦ 0.3%, P ≦ 0.02%, Al ≦ 0.1%, S ≦ 0.01%, N ≦ 0.01%, Ti: 0.05 to 0.35%, and one or more selected from Mo ≦ 0.6% and W ≦ 1.5%, and Mo and W are each independently When it is contained, Mo ≧ 0.1%, W ≧ 0.2%, the balance is made of Fe and inevitable impurities, and Ex. C is 0.015% or less, and Mn ≦ 1.7-30 × Ex. C is satisfied, and 95% or more by volume% is a ferrite structure, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed in the ferrite structure. A high-formability, high-tensile steel plate with excellent strength stability with a tensile strength of 550 MPa or more.
Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184) (1)
(However, the element symbol in the formula (1) indicates mass% of each element.)
Ex.C=C−{Ti−N×(48/14)−S×(48/32)}×(12/48)−Mo×(12/96)−W×(12/184)…(1)
(ただし、(1)式の元素記号は各元素の質量%を示す。) In mass%, C: 0.03 to 0.15%, Mn ≧ 0.2%, Si ≦ 0.3%, P ≦ 0.02%, Al ≦ 0.1%, S ≦ 0.01%, N ≦ 0.01%, Ti: 0.05 to 0.35%, and one or more selected from Mo ≦ 0.6% and W ≦ 1.5%, and Mo and W are each independently When included, Mo ≧ 0.1%, W ≧ 0.2%, and Cr ≦ 0.5%, V ≦ 0.1%, Nb ≦ 0.08%, B ≦ 0.001 %, And the balance consists of Fe and unavoidable impurities, and is represented by the following formula (1). C is 0.015% or less, and Mn ≦ 1.7-30 × Ex. C is satisfied, and 95% or more by volume% is a ferrite structure, and precipitates of less than 10 nm containing Ti and one or more of Mo and W are dispersed in the ferrite structure. A high-formability, high-tensile steel plate with excellent strength stability with a tensile strength of 550 MPa or more.
Ex. C = C- {Ti-N * (48/14) -S * (48/32)} * (12/48) -Mo * (12/96) -W * (12/184) (1)
(However, the element symbol in the formula (1) indicates mass% of each element.)
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