JP3646539B2 - Manufacturing method of hot-dip galvanized high-tensile steel sheet with excellent workability - Google Patents
Manufacturing method of hot-dip galvanized high-tensile steel sheet with excellent workability Download PDFInfo
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Description
【0001】
【発明の属する技術分野】
本発明は、自動車車体などに用いられる、加工性に優れた溶融亜鉛めっき高張力鋼板(合金化したものを含む)の製造方法に関する。
【0002】
【従来の技術】
自動車用の鋼板には、一般に、耐食性と加工性が必要であるため、種々の表面処理鋼板が用いられている。なかでも、溶融亜鉛めっき鋼板は、高度な耐食性を有しているとともに、再結晶焼鈍および亜鉛めっきを同一ラインで処理できる連続溶融亜鉛めっきライン(CGL)により、極めて安価に製造できるという利点を具えている。また、前記亜鉛めっきの後、引き続いて合金化処理を行った溶融亜鉛めっき鋼板は、とりわけ耐食性に優れ、溶接性やプレス成形性にも優れている。
一方、近年、地球環境の改善を目指した燃費向上のための自動車の軽量化が迫られ、また、安全性向上のための衝突時の安全規制の強化が要請されるようになって、溶融亜鉛めっき鋼板にも高強度化(高張力化)が必要になってきた。
【0003】
ところで、高張力鋼板には種々の強化機構を利用したものが開発されているが、とりわけ、自動車の耐衝突特性に優れた鋼板として複合組織鋼板が挙げられる。複合組織鋼板は、フェライト相に、第2相として、主にマルテンサイト相を複合させた鋼板であり、この硬質な第2相を分散させることによって、組織強化による高強度化を図ったものである。
複合組織鋼板の一般的な製造方法は、低炭素材にMnなどの合金元素を添加し、フェライトとオーステナイトの2相領域に加熱したのち、冷却し、オーステナイト相をマルテンサイトに低温変態させるものである。このマルテンサイト変態時に、マルテンサイトの周囲のフェライトに可動転位が導入され、降伏比(YR=降伏強さ/引張強さ)が低くなる。このように降伏比が低い材料では、プレス成形時のしわの発生が抑えられるので、プレス成形に有利である。また、複合組織鋼板には、加工硬化(n値)が高く、均一伸びが高いという利点もある。
【0004】
上述した2相域焼鈍において、オーステナイト相をマルテンサイト相に変態させるためには、合金元素の添加が必要である。例えば、特開昭57−152421号公報には、焼鈍後の冷却速度に応じて合金元素の添加量を規定する技術が提案されている。この開示技術のように、焼き入れ性を向上させるためには、Mn、Mo、Crなどの合金元素を添加する必要がある。
ここで、Moはめっき性への影響が小さいものの、コストアップを招き、多量に添加することは難しい。このため、高強度化を図るための合金添加は、主としてMnあるいはCrを添加することによって対処していた。
【0005】
しかし、このMn、Crは、一般に、焼鈍の過程で鋼板の表面に濃化して (表面濃化層の形成) 、めっき性、とくに溶融亜鉛めっきする際の濡れ性を悪くし不めっきをもたらすことが知られているので、極力低減することが望ましい元素であるといえる。一方、溶融亜鉛めっき後に合金化処理する場合には、通常の連続焼鈍ライン(CAL)の場合に比べて冷却速度が遅くなるので、マルテンサイトを確保するためには、より多くの合金元素の添加が避けられなくなるという側面がある。このため、溶融亜鉛めっき後、合金化処理した溶融亜鉛めっき高張力鋼板で、低降伏比の特性が得られる程度に合金元素を添加すると、他方で不めっきが発生し、外観を問題視する自動車用の部品への適用が困難になるという問題があった。
【0006】
これらの問題に対する従来の方策としては、めっき濡れ性の改善について、例えば、鋼板を連続亜鉛めっきラインに導入するに先立って、電気めっきを行う方法(特開平2−194156号公報)、クラッド法によりSi、Mnなどの含有量の少ない組成の鋼を表層にする方法(特開平3−199363号公報)が提案されている。しかし、これらの方法では、コストがかかり、生産性も悪い工程を新たに経る必要があるなどの問題が生じてしまう。
【0007】
【発明が解決しようとする課題】
上述したように、溶融亜鉛めっき高張力鋼板を製造するに当たって、従来の既知技術は、不めっきの発生、密着性の低下、降伏比の上昇(加工性の低下)などを招き、また、過度の合金添加や新たな附帯設備の増設などに伴って、生産性の低下やコストの上昇をもたらしていた。
本発明は、従来技術が抱えていたこのような問題を解消した溶融亜鉛めっき高張力鋼板を製造するための新規な製造方法を提案することを目的とする。
また、本発明の目的は、とくに不めっきがなく、密着性に優れ、しかも降伏比が低く、良好な加工性を有する溶融亜鉛めっき高張力鋼板の新規な製造方法を提案することにある。
さらに、本発明の他の目的は、具体的な特性として、引張強さが380 〜1000MPa、とりわけ440 〜580 MPaで、降伏比が55%以下を満たし、めっき性がよい溶融亜鉛めっき高張力鋼板の製造方法を提案することにある。
【0008】
【課題を解決するための手段】
発明者らは、めっき性と加工性を両立させるための溶融亜鉛めっき鋼板の製造方法について、鋭意研究した。その結果、合金元素を適正に添加したうえ、鋼板表面における成分濃化を抑制し、かつ、所望の複合組織を得るために適した、熱処理工程を採用することによって、上記の目的を達成できるとの知見を得て本発明を完成するに至った。その要旨構成は以下のとおりである。
【0009】
(1) C:0.005〜0.15wt%、Mn:0.3〜3.0wt%、Mo:0.05〜1.0wt%を含有するめっき母板を、Ac1変態点以上、A c 3 変態点以下の温度で少なくとも1回は焼鈍し、冷却後に、酸洗して表面濃化層を除去し、次いで、Ac1変態点〜Ac3変態点の温度範囲に加熱し、この加熱温度から少なくともめっき浴温度までの温度域を、B含有量に応じて下記(1)式または(2)式で表される臨界冷却速度以上の速度で冷却して、必要に応じて(すなわち、前記冷却においてめっき浴温度未満まで鋼板を冷却した場合は)少なくともめっき浴温度まで加熱し、次いで(前記めっき浴温度までの加熱の有無にかかわらず) 溶融亜鉛めっきを施し、めっき後300℃までの温度域を、B含有量に応じて下記(1)式または(2)式で表される臨界冷却速度以上の速度で冷却することを特徴とする、加工性に優れた溶融亜鉛めっき高張力鋼板の製造方法。
記
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50 ・・・ (1)
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20 ・・・ (2)
ただし、CR:臨界冷却速度(℃/sec)
【0010】
(2) C:0.005〜0.15wt%、Mn:0.3〜3.0wt%、Mo:0.05〜1.0wt%を含有するめっき母板を、Ac1変態点以上、A c 3 変態点以下の温度で少なくとも1回は焼鈍し、冷却後に、酸洗して表面濃化層を除去し、次いで、Ac1変態点〜Ac3変態点の温度範囲に加熱し、この加熱温度から少なくともめっき浴温度までの温度域を、B含有量に応じて下記(1)式または(2)式で表される臨界冷却速度以上の速度で冷却して、必要に応じて(すなわち、前記冷却においてめっき浴温度未満まで鋼板を冷却した場合は)少なくともめっき浴温度まで加熱し、次いで(前記めっき浴温度までの加熱の有無にかかわらず)溶融亜鉛めっきを施し、引き続いて、合金化処理を行い、合金化処理後300℃までの温度域を、B含有量に応じて下記(1)式または(2)式で表される臨界冷却速度以上の速度で冷却することを特徴とする、加工性に優れた溶融亜鉛めっき高張力鋼板の製造方法。
記
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50 ・・・ (1)
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20 ・・・ (2)
ただし、CR:臨界冷却速度(℃/sec)
【0011】
(3) 上記(1)または(2)において、めっき用母板の成分組成が、C:0.005〜0.15wt%、Mn:0.3〜3.0wt%、Mo:0.05〜1.0wt%、を含み、かつSi:0.05〜0.5wt%、Cr:0.05〜1.0wt%、P:0.02〜0.1wt%、B:0.0003〜0.01wt%、Ni:0.05〜1.5wt%、Cu:0.05〜1.5wt%、Nb:0.3wt%以下、Ti:0.3wt%以下、およびV:0.3wt%以下から選ばれるいずれか1種または2種以上を含有し、残部はFeおよび不可避的不純物からなることを特徴とする、加工性に優れた溶融亜鉛めっき高張力鋼板の製造方法。
【0012】
【発明の実施の形態】
まず、本発明の成分組成を上記範囲に限定理由したについて説明する。
C:0.005 〜0.15wt%
Cは、第2相をマルテンサイト化し、また、そのマルテンサイト相の強度を確保するために必要な元素である。C量が0.005 wt%未満では、マルテンサイト化しにくく、複合組織を安定して得ることが困難となる。一方、0.15wt%を超えて添加するとマルテンサイトへの変態温度が低下し、マルテンサイト化しにくくなる。このため、C量は0.005 〜0.15wt%、好ましくは0.02〜0.10wt%とする。
【0013】
Mn:0.3 〜3.0 wt
Mnは、焼き入れ性を向上させる元素として有効な元素であり、安定した複合組織を得るためには、少なくとも0.3 wt%は添加する必要がある。一方、Mn含有量が3.0 wt%を超えると、加工性が低下し、また、めっき性が本発明工程によっても改善できなくなる。このため、Mn量は、0.3 〜3.0 wt%、好ましくは 1.0〜2.4 wt%とする。
【0014】
Mo:0.05〜1.0 wt%
Moは、焼き入れ性を向上させるが、めっき性への悪影響が少ないので、強度確保の上で極めて有用な元素である。このような効果を発揮させるためには、0.05wt%以上の添加が必要である。一方、1.0 wt%を超える添加は、合金化の遅延を招くほか、コスト上昇にもつながるので、Mo添加量は0.05〜1.0 wt%、好ましくは0.10〜0.50wt%の範囲とする。
本発明の鋼板は、上記組成を基本成分として、残部はFeおよび不可避的不純物とすればよい。
【0015】
以上の基本成分に加えて、高張力鋼板のさらなる材質改善をはかるために、焼入性改善元素としてSi, Cr, P, B, Ni, Cuのうち1種以上を、また、局部延性改善元素としてTi, Nb, Vのうち1種以上を、それぞれ添加してよい。
【0016】
Si:0.05〜0.5 wt%
Siは、鋼の強化と強度−伸びバランスに有用な元素である。その効果は0.05wt%以上の添加で得られるが、0.5 wt%を超えて添加すると、めっき性、とくに濡れ性を阻害する。このため、Si量は0.05〜0.5 wt%とする。
【0017】
Cr:0.05〜1.0 wt%
Crは、マルテンサイト化を促進するとともに、マルテンサイトの分布状態を制御し、低降伏比化に有利な元素である。この効果は0.05wt%以上の添加で発現するが、1.0 wt%を超えて添加すると濡れ性を阻害する。よって、Cr量は、0.05〜1.0 wt%の範囲で添加する。
【0018】
P:0.02〜0.1 wt%
Pは、強度向上のほか、伸びやr値の改善に有効な元素である。これらの効果は0.02wt%以上で得られるが、0.1 wt%を超えての添加は、加工性の低下、靭性の低下をもたらすので、0.02〜0.1 wt%の範囲とする。
【0019】
B:0.0003〜0.01wt%
Bは、焼き入れ性を改善するほか、伸びの改善に有効な元素である。この効果は0.0003wt%以上で得られるが、0.01wt%%を超えて添加すると析出による加工性の低下をきたす。よって、Bは0.0003〜0.01wt%の範囲で添加する。
【0020】
Ni:0.05〜1.5 wt%
Niは、焼き入れ性の向上に有効であり、めっき性への悪影響が少ない元素であるが、1.5 wt%を超えて添加すると伸びなどの加工特性を低下させるので、0.05〜1.5 wt%の範囲で添加する。
【0021】
Cu:0.05〜1.5 wt%
Cuは、焼き入れ性を向上させる元素である。この効果を発揮させるためには、0.05wt%以上の添加が必要であるが、1.5 wt%を超えて添加すると、熱間圧延におけるスケール疵の原因になりやすいので、0.05〜1.5 wt%の範囲で添加する。
【0022】
Nb:0.3 wt%以下、Ti:0.3 wt%以下、V:0.3 wt%以下
Nb, Ti, Vは、微細な炭化物をフェライトへ析出させることによりフェライトの強度を上昇させ、伸びフランジ性などの局部延性を向上させるのに有効な元素である。但し、0.3 wt%を超えての添加は析出物が多くなりすぎて伸びの低下をまねくので、0.3 wt%を上限とする。
なお、Nb、TiおよびVの3元素は、同等の効果をもち、0.3 wt%超の添加は伸びの低下をまねくので、合計量 (Nb+Ti+V量) で0.3 wt%以下の範囲で添加するのが望ましい。
【0023】
次に、溶融亜鉛めっき高張力鋼板の製造条件について説明する。
上述した成分を有するスラブを溶製し、熱間圧延を行う。熱間圧延は、オーステナイト域で終了させることが好ましい。その後、表面の酸化スケールを酸洗により除去したのち、冷間圧延を行い、所定の板厚に調整する。これをAc1変態点以上、Ac3変態点以下の2相域温度に加熱(焼鈍)する。焼鈍は連続焼鈍ライン (以下、単に「CAL」と略記する) が好ましいが、バッチ焼鈍でもよい。
上記CALでの加熱の目的は、鋼板表面に成分の濃化を促すとともに、鋼板を複合組織形成条件におき第2相に合金元素を濃化させることによって、めっき性を改善し、降伏比を低下させることにある。なお、表面直下3〜30μm程度の深さに粒界酸化物が少量形成すると、めっき性にはなおよい。このようなCALでの加熱によって得られる効果は、1回の加熱で十分得られるが、複数回繰り返せば一層大きな効果が期待される。焼鈍時間はCALの場合、積算で30秒〜15分が好ましい。
【0024】
図1は、溶融亜鉛めっき鋼板のめっき性と降伏比に及ぼすCAL加熱温度の影響を示したものである。ここに、溶融亜鉛めっき鋼板は、表1の鋼13(Ac1=710 ℃、Ac3=850 ℃)を種々のCAL温度で加熱した後、連続溶融亜鉛めっきライン(以下単に「CGL」と略記)の加熱温度を760 ℃、保持時間を60秒とし、めっき前(CGL加熱終了からめっき浴浸入まで)およびめっき後300 ℃までを、それぞれ15℃/秒、20℃/秒とする速度で冷却して製造したものである。
図1から、めっき性、降伏比とも良好な加熱温度は、Ac1変態点以上の温度範囲であるといえる。しかし、加熱温度がAc3を超えるγ単相域になると、降伏比は高くなり、また、表面濃化に起因してめっき性も悪化する。
これらの結果をもとに、本発明における焼鈍 (CAL) での鋼板の到達温度は、鋼成分によって定まるAc1変態点以上、Ac3変態点以下 (好ましくは950 ℃以下) とする必要がある。なお、CALの雰囲気は、スケールの発生を抑制するために、鋼板に対し還元性とすることが必要であり、一般には、数%のH2 を含むN2 ガスを用いればよい。
【0025】
通常、CALの還元雰囲気では、鉄は酸化されず、Mnなどが外部酸化層として現れる。ただしミクロ的にみると、粒界にそってMnの内部酸化物が存在することがあり、これが酸洗でもとれずに残る。
また、CALで二相組織にすることは、フェライトの粒界にMnなどの元素が濃化した第2相を存在させることになるので、次工程でのCGLのサイクルで、それらの元素の表面への濃化が抑制できる。
【0026】
上記加熱によって、鋼板表面には、鋼中のPが析出し、Si、Mn、Crなどが酸化物として濃化する。これらの濃化成分は、めっき性に対して悪影響を及ぼすので、この表面濃化層を除去する。このとき、内部酸化物は除去せず残すことが好ましいが、酸洗等の通常の工業的除去手段であれば、内部酸化物は除去されずに鋼板の表層部直下に残る。
【0027】
酸洗に続いて、CGLの還元雰囲気の下で加熱(焼鈍)を行う。このときの温度は、めっき濡れ性や合金化速度の点からは650℃以上あればよいが、複合組織の形成を考えてAc1変態点〜A c 3 変態点の2相域温度とする必要がある。CGLにて、Ac1変態点〜Ac3変態点の温度に加熱することによって、合金元素がさらに第2相つまりγ相へと濃化する。そして、この濃化部分は、その後に所定速度で冷却したとき、マルテンサイト相となって複合組織の形成に寄与する。ここでいう合金元素とは、Mn、Moなどの置換型の合金元素であり、前述したCAL加熱やCGLでの加熱の温度域では、比較的拡散しにくい元素をいう。これらの合金元素は、かかる温度域における加熱をCALおよびCGLで繰り返すことによって、より局所的に濃化し、複合組織がより理想的に安定して形成される。ここで、CALでの加熱の役割は、一旦、フェライトと第2相との複合組織となし、CGLにおけるめっき前の加熱(焼鈍)の際に、第2相にさらに合金元素を濃化させることにより、フェライトおよび第2相がより安定して形成されるようにすることにある。このとき、最終製品と同じ複合組織が得られればもちろんよいが、そうでなくとも、少なくとも粒界の3重点付近に合金元素が濃化するため、最終製品での複合組織形成が安定化する。
【0028】
このように、CALおよびCGLの2工程で加熱すれば、内部酸化層の形成により、めっき性の上からも極めて望ましい表面状態が得られる。
すなわち、CALでの加熱、酸洗による表面濃化層の除去を終えると、表面にミクロ的な内部酸化層が形成された鋼板となる。これを、CGLで加熱すると、めっき性に不利なSiやMnなどの合金元素は、粒界にそって表面に濃化しようとするが、これら合金元素は上記内部酸化物にトラップされて表面には移動できなくなる。一方で、鋼板表面では、還元雰囲気により、還元されたFe層が形成され、めっき性に好ましい表面になる。これらの相乗効果によって、めっき性が著しく向上するのである。
【0029】
CGLにおいて加熱した後に、溶融亜鉛めっき、あるいはさらに合金化処理を施して、最終的に複合組織からなる溶融亜鉛めっき高張力鋼板とするためには、CGLでの加熱温度から少なくともめっき浴温度まで、およびめっき浴温度から (あるいは合金化処理温度から) 300 ℃までの各温度域における冷却速度を以下の条件とすることが必要である。
まず、CGLでの加熱温度からめっき浴温度(通常、 550〜450 ℃)までの温度域では、B含有量に応じて下記式で表される臨界冷却速度CR(℃/sec)以上の速度で冷却する。
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)− 2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)− 2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20
上記各式で冷却することによって、合金元素の濃化部はマルテンサイトに変態し、降伏比が低い複合組織鋼板を製造できる。なお、上式は、合金元素の含有量により、冷却過程でのパーライトの晶出曲線が変化するため、パーライトノーズにかからないよう合金元素の量に応じて冷却速度を制御しなければならないことを意味している。
上記の冷却はめっき浴温度で終えて、そのまま溶融亜鉛めっきを施してもよい。また、めっき浴温度未満まで冷却した後、少なくともめっき浴温度まで加熱して溶融亜鉛めっきを施してもよい。
【0030】
また、溶融亜鉛めっきを施した後、めっき温度から (さらに合金化処理を行う場合には、合金化処理温度(通常、 470℃〜Ac1)から) 300 ℃までの温度域をも、上述した方法と同様にして、B含有量に応じて次式で表される臨界冷却速度以上の速度で冷却する。
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)− 2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)− 2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20
冷却速度が上記速度より小さいと、オーステナイト相がマルテンサイトになる前にベイナイト変態してしまい、製品の降伏比が上昇する。
【0031】
【実施例】
以下に、実施例に基づき本発明について説明する。
表1に示す組成の鋼スラブを、1150℃に加熱したのち、仕上げ温度 900〜 850℃で熱問圧延した。この熱延板を酸洗したあと冷間圧延し、この冷延板をめっき用母板とした。これらの母板をCALで加熱(焼鈍)し、CGLにて、めっき前の加熱(焼鈍)を行い、めっきして、溶融亜鉛めっき鋼板とした。また、めっき後さらに合金化処理を行ったものも製造した。これら工程における、CALでの加熱、CGLでの加熱、めっき、合金化などの処理条件を、表2および以下に示す。
【0032】
【表1】
【0033】
【表2】
【0034】
・CAL、CGLでの加熱(焼鈍・めっき前加熱)
雰囲気:5%H2 +N2 ガス(露点−20℃)
・表面濃化層の除去
塩酸酸洗(濃度:5%Hclの水溶液)
温度:60℃
浸漬時間:6秒
・めっき
めっき浴のAl濃度:0.13wt%
浴温:475 ℃
板温:475 ℃
浸漬時間:3秒
目付け量:45g/m2
なお、表2中、No.13 の鋼は、CGLにおいてめっき前加熱後350 ℃まで冷却し、その後475 ℃まで加熱してめっきを施した。他の鋼は、CGLにおいてめっき前加熱後475 ℃まで冷却し、引き続きめっきを施した。
・合金化
処理温度: 470〜550 ℃
合金化後のFe濃度目標:10wt% (X線を使ったオンライン制御を行った)
【0035】
得られた供試鋼板について、引張特性(降伏強さYS,引張強さTS,伸びEl,降伏伸びYEl,降伏比YR)、めっき性(不めっきの程度)およびパウダリング性を調査した。
・めっき性およびパウダリング性の評価方法
不めっき欠陥の判定は、目視により、不めっき欠陥が全くないものを「1」、もっとも不めっきの多いものを「5」とする5段階で評価した。耐パウダリング性は90°曲げ戻しの後、セロテープに付着した亜鉛粉を蛍光X線にて測定した。蛍光X線は、亜鉛粉の亜鉛の蛍光X線を計数管で2分カウントした。セロテープにうっすらと亜鉛粉が付着した状態が2000cps であり、2500cps 以下であれば、自動車などのプレス成形に耐えうるものとなる。
これらの測定結果を併せて表2に示す。なお、めっき層中Fe含有量は、硫酸によりめつき層を溶解し、原子吸光にて測定した。
【0036】
表2から、本発明によって製造した溶融亜鉛めっき高張力鋼板は、いずれも、合金化処理の有無にかかわらず、めっき性、耐パウダリング性が良好であるとともに、降伏比が54.0以下の低い値が得られることがわかる。
【0037】
【発明の効果】
以上説明したように、本発明によれば、表面濃化層の除去、内部酸化層の形成、複合組織の形成が有効に作用して、優れためっき性と耐パウダリング性、低降伏比を共に満たした、溶融亜鉛めっき高張力鋼板を提供することが可能になる。したがって、本発明は、耐食性と加工性が求められる自動車の車体などの品質向上や生産性の向上に寄与するところが極めて大きい。
【図面の簡単な説明】
【図1】溶融亜鉛めっき鋼板のメッキ性と降伏比に及ぼすCAL加熱温度の影響を示すグラフである。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for producing hot-dip galvanized high-tensile steel sheets (including alloyed ones) that are excellent in workability and are used for automobile bodies and the like.
[0002]
[Prior art]
Since steel plates for automobiles generally require corrosion resistance and workability, various surface-treated steel plates are used. Among them, the hot dip galvanized steel sheet has the advantage that it has a high degree of corrosion resistance and can be manufactured at a very low cost by a continuous hot dip galvanizing line (CGL) that can process recrystallization annealing and galvanizing in the same line. It is. Moreover, the hot-dip galvanized steel sheet that has been subsequently alloyed after the galvanization is particularly excellent in corrosion resistance, and is excellent in weldability and press formability.
On the other hand, in recent years, there has been an urgent need to reduce the weight of automobiles to improve fuel efficiency with the aim of improving the global environment, and there has been a demand for strengthening safety regulations at the time of collision to improve safety. High strength (high tension) has also become necessary for plated steel sheets.
[0003]
By the way, what utilized various reinforcement | strengthening mechanisms is developed as a high-tensile steel plate, Especially a composite structure steel plate is mentioned as a steel plate excellent in the collision-resistant characteristic of a motor vehicle. A composite steel sheet is a steel sheet in which a martensite phase is mainly compounded as a second phase with a ferrite phase. By dispersing this hard second phase, high strength is achieved by strengthening the structure. is there.
The general manufacturing method of a composite steel sheet is to add an alloying element such as Mn to a low carbon material, heat it to the two-phase region of ferrite and austenite, cool it, and transform the austenite phase to martensite at a low temperature. is there. During this martensitic transformation, movable dislocations are introduced into the ferrite surrounding the martensite, and the yield ratio (YR = yield strength / tensile strength) is lowered. Such a material having a low yield ratio is advantageous for press molding because wrinkling during press molding can be suppressed. The composite structure steel plate also has the advantages of high work hardening (n value) and high uniform elongation.
[0004]
In the above-described two-phase annealing, an alloy element needs to be added in order to transform the austenite phase into the martensite phase. For example, Japanese Patent Application Laid-Open No. 57-152421 proposes a technique for defining the amount of alloy element added according to the cooling rate after annealing. In order to improve the hardenability as in this disclosed technique, it is necessary to add an alloy element such as Mn, Mo, or Cr.
Here, although Mo has a small influence on the plating property, it causes an increase in cost and is difficult to add in a large amount. For this reason, alloy addition for increasing the strength has been dealt with mainly by adding Mn or Cr.
[0005]
However, this Mn and Cr are generally concentrated on the surface of the steel sheet during the annealing process (formation of a surface enriched layer), resulting in poor plating properties, particularly wettability during hot dip galvanizing, and non-plating. Therefore, it can be said that it is an element that should be reduced as much as possible. On the other hand, in the case of alloying after hot dip galvanization, the cooling rate is slower than in the case of a normal continuous annealing line (CAL), so in order to secure martensite, the addition of more alloy elements There is an aspect that cannot be avoided. For this reason, after hot dip galvanizing, alloyed hot dip galvanized high-strength steel sheets, when alloying elements are added to such an extent that low yield ratio characteristics can be obtained, the other causes non-plating, and automobiles that have an appearance problem There has been a problem that it is difficult to apply to the parts for manufacturing.
[0006]
Conventional measures against these problems include, for example, a method of performing electroplating prior to introducing a steel sheet into a continuous galvanizing line (Japanese Patent Laid-Open No. 2-194156), a cladding method, for improving plating wettability. A method (Japanese Patent Laid-Open No. Hei 3-199363) has been proposed in which a steel having a low content such as Si and Mn is used as a surface layer. However, these methods are problematic in that they are expensive and require a new process with poor productivity.
[0007]
[Problems to be solved by the invention]
As described above, in producing a hot dip galvanized high-tensile steel sheet, the conventional known techniques cause non-plating, decrease in adhesion, increase in yield ratio (decrease in workability), and the like. With the addition of alloys and the addition of new ancillary facilities, productivity was reduced and costs were increased.
An object of this invention is to propose the novel manufacturing method for manufacturing the hot dip galvanized high-tensile steel plate which solved the problem which the prior art had.
Another object of the present invention is to propose a novel method for producing a hot-dip galvanized high-tensile steel sheet having no unplating, excellent adhesion, a low yield ratio, and good workability.
Furthermore, another object of the present invention is to provide, as specific properties, a hot-dip galvanized high-tensile steel sheet having a tensile strength of 380 to 1000 MPa, particularly 440 to 580 MPa, a yield ratio of 55% or less, and good plating properties. This is to propose a manufacturing method.
[0008]
[Means for Solving the Problems]
Inventors earnestly researched about the manufacturing method of the hot dip galvanized steel plate for making plating property and workability compatible. As a result, the above-mentioned object can be achieved by appropriately adding alloying elements, suppressing the concentration of components on the steel sheet surface, and adopting a heat treatment process suitable for obtaining a desired composite structure. As a result, the present invention was completed. The summary composition is as follows.
[0009]
(1) C: 0.005~0.15wt%, Mn: 0.3~3.0wt%, Mo: plating motherboard containing 0.05 to 1.0%, Ac 1 transformation point or more, at least in A c 3 following transformation point temperature once annealed, after cooling, pickling to remove the surface concentrated layer, then heated to a temperature range of Ac 1 transformation point to Ac 3 transformation point, the temperature from the heating temperature to at least the plating bath temperature The zone is cooled at a rate equal to or higher than the critical cooling rate represented by the following formula (1) or (2) according to the B content, and if necessary (that is, the steel plate to below the plating bath temperature in the cooling) (If the sample is cooled) is heated to at least the plating bath temperature, then hot-dip galvanized (with or without heating to the plating bath temperature), and the temperature range up to 300 ° C. after plating depends on the B content And cooling at a speed higher than the critical cooling speed represented by the following formula (1) or (2): A method for manufacturing hot-dip galvanized high-tensile steel sheets with excellent workability.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50 (1)
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20 (2)
However, CR: critical cooling rate (° C / sec)
[0010]
(2) C: 0.005~0.15wt%, Mn: 0.3~3.0wt%, Mo: plating motherboard containing 0.05 to 1.0%, Ac 1 transformation point or more, at least in A c 3 following transformation point temperature once annealed, after cooling, pickling to remove the surface concentrated layer, then heated to a temperature range of Ac 1 transformation point to Ac 3 transformation point, the temperature from the heating temperature to at least the plating bath temperature The zone is cooled at a rate equal to or higher than the critical cooling rate represented by the following formula (1) or (2) according to the B content, and if necessary (that is, the steel plate to below the plating bath temperature in the cooling) (At the time of cooling) to at least the plating bath temperature, then hot-dip galvanized (with or without heating to the plating bath temperature), followed by alloying treatment, after the alloying treatment 300 ° C The temperature range up to the critical cooling rate represented by the following formula (1) or (2) depending on the B content Wherein the cooling in degrees, the production method of the excellent hot-dip galvanized high-strength steel sheet in workability.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50 (1)
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20 (2)
However, CR: critical cooling rate (° C / sec)
[0011]
(3) In the above (1) or (2), the component composition of the plating base plate includes C: 0.005 to 0.15 wt%, Mn: 0.3 to 3.0 wt%, Mo: 0.05 to 1.0 wt%, and Si: 0.05-0.5 wt%, Cr: 0.05-1.0 wt%, P: 0.02-0.1 wt%, B: 0.0003-0.01 wt%, Ni: 0.05-1.5 wt%, Cu: 0.05-1.5 wt%, Nb: 0.3 wt% or less, Ti: 0.3 wt% or less, and V: containing any one or more selected from 0.3 wt% or less, the balance being composed of Fe and inevitable impurities Of hot-dip galvanized high-tensile steel sheet with excellent properties.
[0012]
DETAILED DESCRIPTION OF THE INVENTION
First, the reason why the component composition of the present invention is limited to the above range will be described.
C: 0.005 to 0.15 wt%
C is an element necessary for converting the second phase into martensite and ensuring the strength of the martensite phase. If the amount of C is less than 0.005 wt%, it is difficult to form martensite and it is difficult to stably obtain a composite structure. On the other hand, if it exceeds 0.15 wt%, the transformation temperature to martensite is lowered and it becomes difficult to convert to martensite. For this reason, the C content is 0.005 to 0.15 wt%, preferably 0.02 to 0.10 wt%.
[0013]
Mn: 0.3 to 3.0 wt
Mn is an effective element as an element for improving the hardenability, and it is necessary to add at least 0.3 wt% in order to obtain a stable composite structure. On the other hand, when the Mn content exceeds 3.0 wt%, the workability is lowered, and the plating property cannot be improved even by the process of the present invention. For this reason, the amount of Mn is 0.3 to 3.0 wt%, preferably 1.0 to 2.4 wt%.
[0014]
Mo: 0.05-1.0 wt%
Mo improves the hardenability but has little adverse effect on the plating properties, so it is an extremely useful element for securing strength. In order to exert such an effect, it is necessary to add 0.05 wt% or more. On the other hand, addition exceeding 1.0 wt% causes a delay in alloying and leads to an increase in cost. Therefore, the Mo addition amount is set to 0.05 to 1.0 wt%, preferably 0.10 to 0.50 wt%.
In the steel sheet of the present invention, the above composition may be a basic component, and the balance may be Fe and inevitable impurities.
[0015]
In addition to the above basic components, in order to further improve the material quality of high-strength steel sheets, one or more of Si, Cr, P, B, Ni, and Cu are used as a hardenability improving element, and a local ductility improving element. As above, one or more of Ti, Nb, and V may be added.
[0016]
Si: 0.05-0.5 wt%
Si is an element useful for strengthening steel and strength-elongation balance. The effect can be obtained by addition of 0.05 wt% or more, but if it exceeds 0.5 wt%, the plating property, particularly wettability, is inhibited. For this reason, the Si content is 0.05 to 0.5 wt%.
[0017]
Cr: 0.05-1.0 wt%
Cr is an element that promotes martensite formation and controls the distribution of martensite and is advantageous for lowering the yield ratio. This effect is manifested when 0.05 wt% or more is added, but if it exceeds 1.0 wt%, the wettability is inhibited. Therefore, the Cr amount is added in the range of 0.05 to 1.0 wt%.
[0018]
P: 0.02-0.1 wt%
P is an element effective for improving strength and improving elongation and r value. These effects can be obtained at 0.02 wt% or more. However, addition exceeding 0.1 wt% brings about a decrease in workability and a decrease in toughness, so the range is 0.02 to 0.1 wt%.
[0019]
B: 0.0003-0.01wt%
B is an element effective for improving hardenability and improving elongation. This effect is obtained at 0.0003 wt% or more, but if it exceeds 0.01 wt%, workability is reduced due to precipitation. Therefore, B is added in the range of 0.0003 to 0.01 wt%.
[0020]
Ni: 0.05-1.5 wt%
Ni is an element that is effective in improving hardenability and has little adverse effect on plating properties, but if added over 1.5 wt%, it degrades the processing characteristics such as elongation, so it is in the range of 0.05 to 1.5 wt%. Add in.
[0021]
Cu: 0.05 to 1.5 wt%
Cu is an element that improves hardenability. In order to exert this effect, it is necessary to add 0.05 wt% or more, but if it exceeds 1.5 wt%, it tends to cause scale flaws in hot rolling, so the range of 0.05 to 1.5 wt% Add in.
[0022]
Nb: 0.3 wt% or less, Ti: 0.3 wt% or less, V: 0.3 wt% or less
Nb, Ti and V are effective elements for increasing the strength of ferrite by precipitating fine carbides on ferrite and improving local ductility such as stretch flangeability. However, addition exceeding 0.3 wt% causes an increase in precipitates and decreases elongation, so 0.3 wt% is the upper limit.
The three elements Nb, Ti, and V have the same effect, and addition of more than 0.3 wt% leads to a decrease in elongation. Therefore, the total amount (Nb + Ti + V amount) should be added in the range of 0.3 wt% or less. desirable.
[0023]
Next, manufacturing conditions for the hot-dip galvanized high-tensile steel sheet will be described.
A slab having the components described above is melted and hot rolled. The hot rolling is preferably finished in the austenite region. Then, after removing the oxide scale on the surface by pickling, cold rolling is performed to adjust to a predetermined plate thickness. This is heated (annealed) to a two-phase region temperature not lower than the Ac 1 transformation point and not higher than the Ac 3 transformation point. The annealing is preferably a continuous annealing line (hereinafter simply abbreviated as “CAL”), but may be batch annealing.
The purpose of heating in the CAL is to promote concentration of components on the surface of the steel sheet, and to improve the plating property by concentrating the alloy element in the second phase by placing the steel sheet in the complex structure forming condition, and to improve the yield ratio. It is to reduce. In addition, when a small amount of grain boundary oxide is formed at a depth of about 3 to 30 μm immediately below the surface, the plating property is still good. The effect obtained by such heating with CAL can be sufficiently obtained by one heating, but a larger effect can be expected if it is repeated a plurality of times. In the case of CAL, the annealing time is preferably 30 seconds to 15 minutes.
[0024]
FIG. 1 shows the influence of the CAL heating temperature on the plateability and yield ratio of a hot dip galvanized steel sheet. Here, the hot dip galvanized steel sheet is obtained by heating the steel 13 in Table 1 (Ac 1 = 710 ° C., Ac 3 = 850 ° C.) at various CAL temperatures and then abbreviated as “CGL”. ) Heating temperature of 760 ° C, holding time of 60 seconds, cooling before plating (from the end of CGL heating to entering the plating bath) and 300 ° C after plating at a rate of 15 ° C / second and 20 ° C / second, respectively. It is manufactured.
From FIG. 1, it can be said that the heating temperature with good plating properties and yield ratio is in the temperature range above the Ac 1 transformation point. However, when the heating temperature is in the γ single-phase region exceeding Ac 3 , the yield ratio becomes high, and the plating property also deteriorates due to surface concentration.
Based on these results, the ultimate temperature of the steel sheet during annealing (CAL) in the present invention needs to be not less than the Ac 1 transformation point determined by the steel composition and not more than the Ac 3 transformation point (preferably 950 ° C. or less). . The CAL atmosphere needs to be reducible with respect to the steel sheet in order to suppress the generation of scale, and in general, N 2 gas containing several percent of H 2 may be used.
[0025]
Normally, in a CAL reducing atmosphere, iron is not oxidized and Mn or the like appears as an external oxide layer. However, when viewed microscopically, an internal oxide of Mn may exist along the grain boundary, which remains without being pickled.
In addition, when the CAL has a two-phase structure, a second phase enriched with an element such as Mn is present at the ferrite grain boundary, and therefore the surface of these elements in the CGL cycle in the next step. Concentration to can be suppressed.
[0026]
By the heating, P in the steel is precipitated on the surface of the steel sheet, and Si, Mn, Cr, etc. are concentrated as oxides. Since these thickening components adversely affect the plating property, this surface thickening layer is removed. At this time, it is preferable to leave the internal oxide without removing it, but if it is a normal industrial removal means such as pickling, the internal oxide is not removed but remains directly under the surface layer portion of the steel sheet.
[0027]
Following pickling, heating (annealing) is performed under a reducing atmosphere of CGL. At this time, the temperature may, if 650 ° C. or higher in terms of the plating wettability and alloying rate, needs to be 2-phase region temperature of Ac 1 transformation point ~ A c 3 transformation point consider the formation of a composite structure There is. In CGL, the alloy element is further concentrated to the second phase, that is, the γ phase by heating to a temperature of Ac 1 transformation point to Ac 3 transformation point. Then, when this concentrated portion is subsequently cooled at a predetermined rate, it becomes a martensite phase and contributes to the formation of the composite structure. The alloy element here is a substitutional alloy element such as Mn or Mo, and means an element that is relatively difficult to diffuse in the temperature range of CAL heating or CGL heating described above. These alloy elements are concentrated more locally by repeating heating in such a temperature range by CAL and CGL, and a composite structure is formed more ideally and stably. Here, the role of heating in CAL is once to form a composite structure of ferrite and the second phase, and in heating (annealing) before plating in CGL, the alloy element is further concentrated in the second phase. Thus, the ferrite and the second phase are formed more stably. At this time, it is of course sufficient if the same composite structure as that of the final product is obtained, but even if this is not the case, the alloy element is concentrated at least near the triple point of the grain boundary, so that formation of the composite structure in the final product is stabilized.
[0028]
As described above, when heating is performed in two steps of CAL and CGL, an extremely desirable surface state can be obtained from the viewpoint of plating properties by forming the internal oxide layer.
That is, when the removal of the surface concentrated layer by heating with CAL and pickling is finished, the steel sheet has a microscopic internal oxide layer formed on the surface. When this is heated with CGL, alloy elements such as Si and Mn, which are disadvantageous in plating properties, try to concentrate on the surface along the grain boundaries. However, these alloy elements are trapped by the internal oxide and are deposited on the surface. Can no longer move. On the other hand, on the steel sheet surface, a reduced Fe layer is formed in a reducing atmosphere, and the surface becomes preferable for plating properties. These synergistic effects significantly improve the plating properties.
[0029]
After heating in CGL, hot-dip galvanization, or further alloying treatment, to finally form a hot-dip galvanized high-tensile steel sheet composed of a composite structure, from the heating temperature in CGL to at least the plating bath temperature, In addition, the cooling rate in each temperature range from the plating bath temperature (or from the alloying treatment temperature) to 300 ° C. must be as follows.
First, in the temperature range from the heating temperature in CGL to the plating bath temperature (usually 550 to 450 ° C.), at a rate equal to or higher than the critical cooling rate CR (° C./sec) represented by the following formula according to the B content. Cooling.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20
By cooling with the above formulas, the concentrated portion of the alloy element is transformed into martensite, and a composite structure steel plate having a low yield ratio can be manufactured. The above equation means that the crystallization curve of pearlite in the cooling process changes depending on the content of the alloy element, so the cooling rate must be controlled according to the amount of alloy element so as not to cause pearlite nose. doing.
The above cooling may be finished at the plating bath temperature and hot dip galvanizing may be performed as it is. Further, after cooling to below the plating bath temperature, the hot dip galvanizing may be performed by heating to at least the plating bath temperature.
[0030]
In addition, after the hot dip galvanization, the temperature range from the plating temperature to 300 ° C. (when further alloying treatment is performed, from the alloying treatment temperature (usually 470 ° C. to Ac 1 )) is also described above. In the same manner as in the method, cooling is performed at a speed equal to or higher than the critical cooling speed represented by the following formula according to the B content.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20
When the cooling rate is lower than the above rate, the austenite phase is transformed into bainite before becoming martensite, and the yield ratio of the product increases.
[0031]
【Example】
Hereinafter, the present invention will be described based on examples.
A steel slab having the composition shown in Table 1 was heated to 1150 ° C. and then hot rolled at a finishing temperature of 900 to 850 ° C. The hot-rolled sheet was pickled and cold-rolled, and the cold-rolled sheet was used as a plating base plate. These base plates were heated (annealed) with CAL, heated (annealed) before plating with CGL, and plated to obtain hot-dip galvanized steel sheets. Moreover, the thing which performed the alloying process after plating was also manufactured. Table 2 and the following show the processing conditions such as heating with CAL, heating with CGL, plating, and alloying in these steps.
[0032]
[Table 1]
[0033]
[Table 2]
[0034]
-Heating with CAL and CGL (annealing / heating before plating)
Atmosphere: 5% H 2 + N 2 gas (dew point -20 ° C)
・ Removal of concentrated surface layer Hydrochloric acid pickling (concentration: 5% Hcl aqueous solution)
Temperature: 60 ° C
Immersion time: 6 seconds ・ Al concentration in plating bath: 0.13wt%
Bath temperature: 475 ℃
Plate temperature: 475 ℃
Immersion time: 3 seconds Weight per unit area: 45 g / m 2
In Table 2, the No. 13 steel was cooled to 350 ° C. after heating before plating in CGL, and then heated to 475 ° C. for plating. The other steels were cooled to 475 ° C. after pre-plating heating in CGL and subsequently plated.
-Alloying temperature: 470-550 ° C
Fe concentration target after alloying: 10wt% (On-line control using X-ray was performed)
[0035]
The obtained test steel sheet was examined for tensile properties (yield strength YS, tensile strength TS, elongation El, yield elongation YEl, yield ratio YR), plating property (degree of non-plating) and powdering property.
-Evaluation method of plating property and powdering property The determination of the non-plating defect was visually evaluated in five stages, with 1 indicating that there was no non-plating defect and 5 indicating the most non-plating defect. The powdering resistance was measured by fluorescent X-rays after 90 ° bending back and zinc powder adhering to the cello tape. For fluorescent X-rays, zinc fluorescent X-rays of zinc powder were counted for 2 minutes with a counter. The state where the zinc powder is slightly adhered to the cellophane is 2000 cps, and if it is 2500 cps or less, it can withstand press forming of an automobile or the like.
These measurement results are shown together in Table 2. The Fe content in the plating layer was measured by atomic absorption by dissolving the plating layer with sulfuric acid.
[0036]
From Table 2, the hot-dip galvanized high-strength steel sheets produced by the present invention all have good plating properties and powdering resistance regardless of the presence of alloying treatment, and the yield ratio is a low value of 54.0 or less. It can be seen that
[0037]
【The invention's effect】
As described above, according to the present invention, the removal of the surface concentrated layer, the formation of the internal oxide layer, and the formation of the composite structure function effectively, and excellent plating properties, powdering resistance, and a low yield ratio are achieved. It is possible to provide a hot-dip galvanized high-tensile steel sheet that satisfies both requirements. Therefore, the present invention greatly contributes to improving the quality and productivity of automobile bodies that require corrosion resistance and workability.
[Brief description of the drawings]
FIG. 1 is a graph showing the influence of CAL heating temperature on the plateability and yield ratio of a hot-dip galvanized steel sheet.
Claims (3)
記
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50 ・・・ (1)
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20 ・・・ (2)
ただし、CR:臨界冷却速度(℃/sec)C: 0.005~0.15wt%, Mn: 0.3~3.0wt %, Mo: the 0.05 to 1.0% plating motherboard containing, Ac 1 transformation point or more, at least once with A c 3 following transformation point temperature Is annealed, cooled, pickled to remove the surface concentrated layer, and then heated to a temperature range of Ac 1 transformation point to Ac 3 transformation point, and a temperature range from this heating temperature to at least the plating bath temperature, Depending on the B content, it is cooled at a rate equal to or higher than the critical cooling rate represented by the following formula (1) or (2), heated to at least the plating bath temperature as required, and then hot dip galvanized. Excellent in workability, characterized by cooling the temperature range up to 300 ° C after plating at a rate equal to or higher than the critical cooling rate represented by the following formula (1) or (2) according to the B content Manufacturing method of hot dip galvanized high strength steel sheet.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50 (1)
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20 (2)
However, CR: critical cooling rate (° C / sec)
記
B≦0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.50 ・・・ (1)
B>0.0006wt%のとき、
log CR=−3.50(Mowt%)−1.20(Mnwt%)−0.16(Siwt%)−2.0(Crwt%)−0.08(Niwt%+Cuwt%)−0.32(Pwt%)+3.20 ・・・ (2)
ただし、CR:臨界冷却速度(℃/sec)C: 0.005~0.15wt%, Mn: 0.3~3.0wt %, Mo: the 0.05 to 1.0% plating motherboard containing, Ac 1 transformation point or more, at least once with A c 3 following transformation point temperature Is annealed, pickled to remove the surface concentrated layer after cooling, and then heated to a temperature range from Ac 1 transformation point to Ac 3 transformation point, and a temperature range from this heating temperature to at least the plating bath temperature. Depending on the B content, cool at a rate equal to or higher than the critical cooling rate represented by the following formula (1) or (2), and if necessary, heat to at least the plating bath temperature, and then apply hot dip galvanization. Subsequently, the alloying treatment is performed, and the temperature range up to 300 ° C. after the alloying treatment is cooled at a speed equal to or higher than the critical cooling rate represented by the following formula (1) or (2) according to the B content. A method for producing a hot-dip galvanized high-tensile steel sheet having excellent workability.
When B ≦ 0.0006wt%,
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.50 (1)
When B> 0.0006wt%
log CR = −3.50 (Mowt%) − 1.20 (Mnwt%) − 0.16 (Siwt%) − 2.0 (Crwt%) − 0.08 (Niwt% + Cuwt%) − 0.32 (Pwt%) + 3.20 (2)
However, CR: critical cooling rate (° C / sec)
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