JP2009280899A - METHOD FOR PRODUCING HIGH-STRENGTH HOT-ROLLED STEEL SHEET HAVING 780 MPa OR MORE OF TENSILE STRENGTH - Google Patents
METHOD FOR PRODUCING HIGH-STRENGTH HOT-ROLLED STEEL SHEET HAVING 780 MPa OR MORE OF TENSILE STRENGTH Download PDFInfo
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本発明は、自動車の足回り部品などに適した板厚8mm以下の高強度熱延鋼板、特に、加工後の伸びフランジ性に優れ、かつ鋼板内材質変動の小さい780MPa以上の引張強度TSを有する高強度熱延鋼板の製造方法に関する。 The present invention is a high-strength hot-rolled steel sheet having a thickness of 8 mm or less suitable for automobile undercarriage parts, etc., and in particular, has a tensile strength TS of 780 MPa or more which is excellent in stretch flangeability after processing and small in material variation in the steel sheet The present invention relates to a method for producing a high-strength hot-rolled steel sheet.
近年、環境問題に対する関心が高まるなか、自動車用鋼板には、軽量化による燃費向上を目的に一層の高強度-薄肉化が要求されている。特に、自動車の足回り部品などには、現状多用されている440MPa程度のTSを有する高強度熱延鋼板に代わって、590MPa以上さらには780MPa以上のTSを有する高強度熱延鋼板の使用が増大する傾向にある。 In recent years, with increasing interest in environmental issues, steel sheets for automobiles are required to have higher strength and thinner wall thickness for the purpose of improving fuel efficiency through weight reduction. In particular, the use of high-strength hot-rolled steel sheets with TS of 590 MPa or higher and even 780 MPa or higher is increasing for undercarriage parts of automobiles in place of high-strength hot-rolled steel sheets with TS of about 440 MPa, which are widely used at present. Tend to.
自動車の足回り部品を製造するには、厳しい伸びフランジ加工がともなう場合が多いので、こうした高強度熱延鋼板には、優れた伸びフランジ性、特に、近年のプレス技術の進歩により、伸びフランジ加工はドロー(絞りおよび張り出し)→トリム(穴抜き)→リストライク(穴広げ)のような工程で行われる場合が増加しているため、優れた加工後の伸びフランジ性が必要である。 Manufacturing automobile undercarriage parts often involves severe stretch flange processing, so these high-strength hot-rolled steel sheets have excellent stretch flangeability, especially due to recent advances in press technology. Since there is an increasing number of cases of drawing (drawing and overhanging) → trimming (hole punching) → wrist like (hole expanding), an excellent stretch flangeability after processing is required.
従来より、780MPa以上のTSを有し、かつ伸びフランジ性に優れた高強度熱延鋼板がいくつか提案されている。例えば、特許文献1には、質量%で、C:0.06〜0.12%、Si:0.5〜1.6%、Mn:1.5〜2.2%、S:0.0050%以下、Ti:0.03〜0.08%を含有し、残部Feおよび不可避的不純物からなり、体積率3%以上の残留オーステナイトおよび微細ベイナイトからなる組織を有する伸びフランジ性および延性に優れた高強度熱延鋼板およびその製造方法が開示されている。特許文献2には、質量%で、C:0.06〜0.15%、Si:1.2%以下、Mn:0.5〜1.6%、P:0.04%以下、S:0.005%以下、Al:0.05%以下、Ti:0.03〜0.20%を含有し、残部Feおよび不可避的不純物からなる成分組成を有するとともに、体積率50〜90%がフェライトで、かつ残部が実質的にベイナイトであって、フェライトとベイナイトの体積率の合計が95%以上であり、フェライト中にはTiを含む析出物が析出し、該析出物の平均直径が20nm以下である組織を有し、かつ、鋼中のTi量の80%以上が析出している伸び特性、伸びフランジ特性および引張疲労特性に優れた高強度熱延鋼板およびその製造方法が開示されている。特許文献3には、質量%で、C:0.05〜0.15%、Si:0.5〜2.5%、Mn:0.5〜2.5%、P:0.02%以下、S:0.01%以下、Al:0.005〜0.1%、Ti:0.3%以下を含み、1.6%≦Si+Mn≦5.0%を満たし、不純物としてB:0.0005%以下に制限し、残部Feおよび不可避的不純物であり、ポリゴナルフェライト占有率VPF(%)とポリゴナルフェライト粒径dPF(μm)の比VPF/dPFが20以上、残留オーステナイトの占有率が5%以上で、残部がベイナイトからなり、強度-延性バランスが20000MPa・%以上で、かつその鋼板内変動が3000MPa・%未満である材質変動の小さい加工性に優れた熱延高強度鋼板およびその製造方法が開示されている。 Conventionally, several high strength hot rolled steel sheets having a TS of 780 MPa or more and excellent stretch flangeability have been proposed. For example, Patent Document 1 contains, in mass%, C: 0.06 to 0.12%, Si: 0.5 to 1.6%, Mn: 1.5 to 2.2%, S: 0.0050% or less, Ti: 0.03 to 0.08%, and the balance A high-strength hot-rolled steel sheet having a structure composed of Fe and inevitable impurities and having a structure composed of retained austenite and fine bainite having a volume ratio of 3% or more and excellent in stretch flangeability and ductility and a method for producing the same are disclosed. Patent Document 2 includes, in mass%, C: 0.06 to 0.15%, Si: 1.2% or less, Mn: 0.5 to 1.6%, P: 0.04% or less, S: 0.005% or less, Al: 0.05% or less, Ti: It contains 0.03 to 0.20%, has a component composition consisting of the balance Fe and unavoidable impurities, a volume ratio of 50 to 90% is ferrite, and the balance is substantially bainite, and the volume ratio of ferrite and bainite The total is 95% or more, precipitates containing Ti are precipitated in the ferrite, the precipitate has an average diameter of 20 nm or less, and 80% or more of the amount of Ti in the steel is precipitated. A high-strength hot-rolled steel sheet having excellent elongation characteristics, stretch flange characteristics and tensile fatigue characteristics, and a method for producing the same are disclosed. Patent Document 3 includes mass%, C: 0.05 to 0.15%, Si: 0.5 to 2.5%, Mn: 0.5 to 2.5%, P: 0.02% or less, S: 0.01% or less, Al: 0.005 to 0.1%, Ti: Including 0.3% or less, 1.6% ≦ Si + Mn ≦ 5.0%, B: 0.0005% or less as impurities, balance Fe and inevitable impurities, Polygonal ferrite occupancy V PF (%) And the ratio of polygonal ferrite particle diameter d PF (μm) V PF / d PF is 20 or more, the retained austenite occupancy is 5% or more, the balance is bainite, the strength-ductility balance is 20000 MPa ・% or more, Also disclosed is a hot-rolled high-strength steel sheet excellent in workability with a small material fluctuation, whose fluctuation in the steel sheet is less than 3000 MPa ·%, and a manufacturing method thereof.
しかしながら、特許文献1に記載の鋼板では、加工を受けるとマルテンサイトに変態する残留オーステナイトを3%以上有する組織であるため、優れた加工後の伸びフランジ性が得られない。特許文献2や3に記載の鋼板では、必ずしも鋼板内材質変動を小さく抑制できず、硬質部では加工後の伸びフランジ性が大きく劣化するという問題がある。 However, since the steel sheet described in Patent Document 1 has a structure having 3% or more of retained austenite that transforms into martensite when subjected to processing, excellent stretch flangeability after processing cannot be obtained. In the steel sheets described in Patent Documents 2 and 3, the material fluctuation in the steel sheet cannot be suppressed to a small level, and there is a problem that the stretch flangeability after processing is greatly deteriorated in the hard part.
本発明は、このような問題を解決するためになされたもので、加工後の伸びフランジ性に優れ、かつ鋼板内材質変動を安定して小さくできる780MPa以上のTSを有する高強度熱延鋼板の製造方法を提供することを目的とする。 The present invention has been made to solve such problems, and is a high-strength hot-rolled steel sheet having a TS of 780 MPa or more, which has excellent stretch flangeability after processing and can stably reduce material fluctuations in the steel sheet. An object is to provide a manufacturing method.
本発明者らは、加工後の伸びフランジ性に優れ、かつ鋼板内材質変動を安定して小さくできる板厚1〜8mm程度で780MPa以上のTSを有する高強度熱延鋼板の製造方法について検討を重ねた結果、以下のことを見出した。 The present inventors examined a method for producing a high-strength hot-rolled steel sheet having a TS of 780 MPa or more with a sheet thickness of about 1 to 8 mm, which has excellent stretch flangeability after processing and can stably reduce the material variation in the steel sheet. As a result of overlapping, the following was found.
i) 780MPa以上のTSで優れた加工後の伸びフランジ性を得るには、熱間圧延後の冷却条件を制御することにより、フェライトのマトリックス中にベイナイトを適量分散させることが効果的である。 i) In order to obtain excellent stretch flangeability after processing with a TS of 780 MPa or more, it is effective to disperse an appropriate amount of bainite in the ferrite matrix by controlling the cooling conditions after hot rolling.
ii) 鋼板内材質変動を安定して小さくするには、鋼板の水冷時に膜沸騰冷却と核沸騰冷却が共存する遷移沸騰冷却となる500℃以下の温度域を120℃/秒以上の冷却速度で核沸騰冷却となる条件で水冷することが効果的である。 ii) In order to stably reduce the material fluctuation in the steel plate, the temperature range of 500 ° C or lower, which is transition boiling cooling in which film boiling cooling and nucleate boiling cooling coexist during water cooling of the steel plate, is set at a cooling rate of 120 ° C / second or more. It is effective to perform water cooling under the condition of nucleate boiling cooling.
本発明は、このような知見に基づいて完成されたもので、質量%で、C:0.04〜0.15%、Si:0.3〜1.5%、Mn:0.5〜2.0%、P:0.06%以下、S:0.005%以下、Al:0.10%以下、Ti:0.07〜0.20%を含み、残部がFeおよび不可避的不純物からなる成分組成を有する鋼片を、1150〜1300℃の加熱温度で加熱し、800〜1000℃の仕上温度で熱間圧延後、30℃/秒以上の冷却速度で600〜750℃の冷却停止温度まで一次冷却し、0.2〜10秒間冷却を停止後、55℃/秒以上の冷却速度で二次冷却するに際し、少なくとも500℃以下の温度域を120℃/秒以上の冷却速度で核沸騰冷却となる条件で冷却し、350〜500℃の巻取温度で巻取ることを特徴とする780MPa以上の引張強度を有する高強度熱延鋼板の製造方法を提供する。 The present invention has been completed based on such findings, and in mass%, C: 0.04 to 0.15%, Si: 0.3 to 1.5%, Mn: 0.5 to 2.0%, P: 0.06% or less, S: A steel slab having a component composition comprising 0.005% or less, Al: 0.10% or less, Ti: 0.07 to 0.20%, and the balance consisting of Fe and inevitable impurities, is heated at a heating temperature of 1150 to 1300 ° C., 800 to 1000 After hot rolling at a finishing temperature of ℃, primary cooling to a cooling stop temperature of 600 to 750 ℃ at a cooling rate of 30 ℃ / second or more, after stopping cooling for 0.2 to 10 seconds, at a cooling rate of 55 ℃ / second or more During secondary cooling, the temperature range of at least 500 ° C. is cooled at a cooling rate of 120 ° C./second or more under the condition of nucleate boiling cooling, and the coil is wound at a coiling temperature of 350 to 500 ° C. A method for producing a high-strength hot-rolled steel sheet having the above tensile strength is provided.
本発明の製造方法では、一次冷却後の冷却停止時間を1〜10秒間にしたり、少なくとも500℃以下の温度域を250℃/秒以上の冷却速度で核沸騰冷却となる条件で冷却することが好ましい。また、前記鋼片が、さらに、質量%で、Cr:0.1〜0.8%、Nb:0.005〜0.1%、V:0.005〜0.2%、W:0.005〜0.2%、Mo:0.01〜0.3%のうちから選ばれた1種または2種以上を含む成分組成を有する鋼片を用いることが好ましい。 In the production method of the present invention, the cooling stop time after the primary cooling can be set to 1 to 10 seconds, or at least a temperature range of 500 ° C. or lower can be cooled at a cooling rate of 250 ° C./second or more under conditions of nucleate boiling cooling. preferable. Further, the steel slab further comprises, in mass%, Cr: 0.1 to 0.8%, Nb: 0.005 to 0.1%, V: 0.005 to 0.2%, W: 0.005 to 0.2%, Mo: 0.01 to 0.3%. It is preferable to use a steel slab having a component composition containing one or more selected ones.
本発明の製造方法により、加工後の伸びフランジ性に優れ、かつ鋼板内材質変動を安定して小さくできる780MPa以上のTSを有する高強度熱延鋼板を製造可能になった。本発明の製造方法で製造された高強度熱延鋼板は、特に、自動車の足回り部品に好適である。 The production method of the present invention makes it possible to produce a high-strength hot-rolled steel sheet having a TS of 780 MPa or more that is excellent in stretch flangeability after processing and that can stably reduce the material fluctuation in the steel sheet. The high-strength hot-rolled steel sheet produced by the production method of the present invention is particularly suitable for automobile underbody parts.
以下に、本発明の高強度熱延鋼板の製造方法の詳細について説明する。なお、各成分元素の含有量を表す「%」は、特に断らない限り「質量%」を意味する。 Below, the detail of the manufacturing method of the high intensity | strength hot-rolled steel plate of this invention is demonstrated. Note that “%” representing the content of each component element means “% by mass” unless otherwise specified.
1) 成分組成
C:0.04〜0.15%
Cは、ベイナイトを生成させ必要な強度を確保するのに必要な元素である。780MPa以上のTSを得るためにはC量を0.04％以上とする必要があるが、0.15％を超えると加工後の伸びフランジ性が低下する。したがって、C量は0.04〜0.15％、好ましくは0.05〜0.10%、より好ましくは0.05〜0.09%とする。
1) Component composition
C: 0.04-0.15%
C is an element necessary for producing bainite and ensuring the necessary strength. In order to obtain a TS of 780 MPa or more, the C amount needs to be 0.04% or more. However, if it exceeds 0.15%, the stretch flangeability after processing is deteriorated. Therefore, the C content is 0.04 to 0.15%, preferably 0.05 to 0.10%, more preferably 0.05 to 0.09%.
Si:0.3〜1.5%
Siは、固溶強化により強度を上昇させるのに必要な元素である。Si量が0.3%未満では780MPa以上のTSを得るために高価な合金元素の添加量を増やす必要がある。一方、1.5%を超えると表面性状の低下を招き疲労特性を劣化させる。したがって、Si量は0.3〜1.5%、好ましくは0.3〜1.2%とする。
Si: 0.3-1.5%
Si is an element necessary for increasing the strength by solid solution strengthening. If the Si content is less than 0.3%, it is necessary to increase the amount of the expensive alloy element added in order to obtain TS of 780 MPa or more. On the other hand, if it exceeds 1.5%, the surface properties are lowered and the fatigue properties are deteriorated. Therefore, the Si content is 0.3 to 1.5%, preferably 0.3 to 1.2%.
Mn:0.5〜2.0%
Mnは、固溶強化およびベイナイト生成に有効な元素である。780MPa以上のTSを得るためにはMn量を0.5%以上とする必要があるが、2.0%を超えると溶接性が低下する。したがって、Mn量は0.5〜2.0%、好ましくは0.8〜0.18%とする。
Mn: 0.5-2.0%
Mn is an element effective for solid solution strengthening and bainite formation. In order to obtain a TS of 780 MPa or more, the Mn content needs to be 0.5% or more, but if it exceeds 2.0%, the weldability deteriorates. Therefore, the Mn content is 0.5 to 2.0%, preferably 0.8 to 0.18%.
P:0.06%以下
P量が0.06%を超えると偏析による加工後の伸びフランジ性の低下を招く。したがって、Pは0.06%以下、好ましくは0.03%以下とする。なお、Pは固溶強化に有効な元素でもあり、この効果を得る上では0.005%以上含有していることが好ましい。
P: 0.06% or less
If the P content exceeds 0.06%, the stretch flangeability after processing is reduced due to segregation. Therefore, P is 0.06% or less, preferably 0.03% or less. Note that P is also an element effective for solid solution strengthening. In order to obtain this effect, P is preferably contained in an amount of 0.005% or more.
S:0.005%以下
Sは、MnおよびTiと硫化物を形成して加工後の伸びフランジ性を低下させるとともに、高強度化に有効なMnやTi量の低減を招く。したがって、S量は0.005%以下とし、より好ましくは0.003%以下とする。なお、S量は極力低減することが好ましい。
S: 0.005% or less
S forms sulfides with Mn and Ti to reduce the stretch flangeability after processing, and causes a reduction in the amount of Mn and Ti effective for increasing the strength. Therefore, the S content is 0.005% or less, more preferably 0.003% or less. Note that the amount of S is preferably reduced as much as possible.
Al:0.10%以下
Alは、鋼の脱酸剤として重要な元素であるが、鋼中のAl量が0.10%を超えると鋼板表面性状の低下を招く。したがって、Al量は0.10%以下、好ましくは0.06%以下とする。なお、脱酸効果を十分に確保する上では、Al量は0.005%以上とすることが好ましい。
Al: 0.10% or less
Al is an important element as a deoxidizer for steel, but if the Al content in the steel exceeds 0.10%, the surface properties of the steel sheet are lowered. Therefore, the Al content is 0.10% or less, preferably 0.06% or less. In order to secure a sufficient deoxidation effect, the Al content is preferably 0.005% or more.
Ti:0.07〜0.20%
Tiは、その一部がCと結合し微細な炭化物を形成し、強度上昇や溶接時のHAZ(熱影響部)軟化防止に寄与する元素である。こうした効果を得るにはTi量を0.07%以上とする必要があるが、0.20%を超えると加工後の伸びフランジ性の低下を招く。したがって、Ti量は0.07〜0.20%、好ましくは0.10%超え0.15%以下とする。
Ti: 0.07 ~ 0.20%
A part of Ti is an element that combines with C to form fine carbides and contributes to strength enhancement and prevention of HAZ (heat affected zone) softening during welding. In order to obtain such an effect, the Ti amount needs to be 0.07% or more. However, if it exceeds 0.20%, the stretch flangeability after processing is lowered. Therefore, the Ti content is 0.07 to 0.20%, preferably more than 0.10% and 0.15% or less.
残部はFeおよび不可避的不純物であるが、以下の理由により、さらに、質量%で、Cr:0.1〜0.8%、Nb:0.005〜0.1%、V:0.005〜0.2%、W:0.005〜0.2%、Mo:0.01〜0.3%のうちから選ばれた少なくとも1種を含有させることが好ましい。 The balance is Fe and inevitable impurities, but for the following reasons, in addition, in mass%, Cr: 0.1 to 0.8%, Nb: 0.005 to 0.1%, V: 0.005 to 0.2%, W: 0.005 to 0.2%, Mo: It is preferable to contain at least one selected from 0.01 to 0.3%.
Cr:0.1〜0.8%
Crは、焼入れ性向上に効果的な元素である。Cr量が0.1%未満ではその効果が小さく、0.8%を超えると加工後の伸びフランジ性の低下を招く。したがって、Cr量は0.1〜0.8%、好ましくは0.3〜0.7%とする。
Cr: 0.1-0.8%
Cr is an element effective for improving hardenability. If the Cr content is less than 0.1%, the effect is small, and if it exceeds 0.8%, the stretch flangeability after processing is lowered. Therefore, the Cr content is 0.1 to 0.8%, preferably 0.3 to 0.7%.
Nb:0.005〜0.1%、V:0.005〜0.2%、W:0.005〜0.2%、Mo:0.01〜0.3%
Nb、V、WおよびMoは、いずれもCと結合し微細な炭化物を形成して強度上昇に寄与する元素である。しかしながら、Nb、V、W量が各々0.005%未満、あるいはMo量が0.01%未満では炭化物の生成量が少なく強度上昇が不十分となり、Nb量が0.1%超え、VおよびW量は各々0.2%超え、あるいはMo量が0.3%超えると加工後の伸びフランジ性の低下を招く。したがって、Nb量は0.005〜0.1%、V量は0.005〜0.2%、W量は0.005〜0.2%、Mo量は0.01〜0.3%とする。
Nb: 0.005-0.1%, V: 0.005-0.2%, W: 0.005-0.2%, Mo: 0.01-0.3%
Nb, V, W, and Mo are all elements that contribute to an increase in strength by combining with C to form fine carbides. However, if the Nb, V, and W amounts are each less than 0.005%, or the Mo amount is less than 0.01%, the amount of carbide generated is small and the strength increase is insufficient, the Nb amount exceeds 0.1%, and the V and W amounts are each 0.2%. Exceeding or if the Mo content exceeds 0.3%, the stretch flangeability after processing is lowered. Therefore, the Nb amount is 0.005 to 0.1%, the V amount is 0.005 to 0.2%, the W amount is 0.005 to 0.2%, and the Mo amount is 0.01 to 0.3%.
なお、本発明の作用効果に害をおよぼさない微量元素として、Cu、Ni、Cr、Sn、Pb、Sbを各々0.1%以下の範囲で含有してもよい。 In addition, Cu, Ni, Cr, Sn, Pb, and Sb may be contained in a range of 0.1% or less as trace elements that do not adversely affect the operational effects of the present invention.
2) 製造条件
熱間圧延前の加熱温度:1150〜1300℃
圧延荷重の低減および良好な表面性状の確保の観点から、加熱温度は1150℃以上とする必要がある。また、熱間圧延前にTiの炭化物、あるいはさらにNb、V、ＷおよびMoを添加した場合は、これらの炭化物を溶解させる上でも、1150℃以上の加熱が必要である。一方、加熱温度が1300℃を超えるとオーステナイト粒が粗大化して圧延後の冷却時にフェライト変態が遅延されるため、ベイナイトの生成量が過剰になり加工後の伸びフランジ性が低下する。したがって、加熱温度は1150〜1300℃とする。
2) Manufacturing conditions Heating temperature before hot rolling: 1150 ~ 1300 ℃
From the viewpoint of reducing the rolling load and ensuring good surface properties, the heating temperature needs to be 1150 ° C. or higher. In addition, when Ti carbide, or Nb, V, W, and Mo are added before hot rolling, heating at 1150 ° C. or higher is necessary to dissolve these carbides. On the other hand, when the heating temperature exceeds 1300 ° C., the austenite grains become coarse and the ferrite transformation is delayed during cooling after rolling, so that the amount of bainite generated becomes excessive, and the stretch flangeability after processing decreases. Accordingly, the heating temperature is 1150-1300 ° C.
熱間圧延の仕上温度:800〜1000℃
仕上温度が800℃未満では圧延荷重の増大による表面欠陥の増加を招く。また、仕上温度が800℃未満ではフェライトとオーステナイトの二相域圧延になり780MPa以上のTSが得られなくなる場合もある。一方、仕上温度が1000℃を超えるとオーステナイト粒が微細化されないため、加工後の伸びフランジ性が低下する。したがって、仕上温度は800〜1000℃、好ましくは820〜950℃とする。
Hot rolling finishing temperature: 800 ~ 1000 ℃
If the finishing temperature is less than 800 ° C., surface defects increase due to an increase in rolling load. Further, if the finishing temperature is less than 800 ° C., the two-phase rolling of ferrite and austenite may occur, and TS of 780 MPa or more may not be obtained. On the other hand, if the finishing temperature exceeds 1000 ° C., the austenite grains are not refined, and the stretch flangeability after processing is lowered. Therefore, the finishing temperature is 800 to 1000 ° C, preferably 820 to 950 ° C.
熱間圧延後の一次冷却条件:冷却速度30℃/秒以上、冷却停止温度600〜750℃
熱間圧延後の一次冷却の冷却速度が30℃/秒未満では高温域からフェライト変態が開始され、ベイナイト生成が困難となり、780MPa以上のTSが得られない。したがって、一次冷却の冷却速度は30℃/秒以上、好ましくは60℃/秒以上として強制冷却する必要がある。なお、ここで、冷却速度は熱間圧延後冷却停止までの一次冷却中の平均冷却速度である。一次冷却の冷却速度の上限は、次に延べる冷却停止温度の精度が確保されれば規制されないが、現状の冷却技術を考慮すると700℃/秒以下が好ましい。なお、一次冷却の方法は、特に限定する必要はなく、例えば、公知のラミナー冷却による水冷を利用できる。
Primary cooling conditions after hot rolling: Cooling rate of 30 ° C / second or more, cooling stop temperature of 600-750 ° C
If the cooling rate of the primary cooling after hot rolling is less than 30 ° C./sec, ferrite transformation starts from a high temperature range, bainite formation becomes difficult, and TS of 780 MPa or more cannot be obtained. Therefore, it is necessary to perform forced cooling at a cooling rate of primary cooling of 30 ° C./second or more, preferably 60 ° C./second or more. Here, the cooling rate is an average cooling rate during primary cooling after hot rolling until cooling is stopped. The upper limit of the cooling rate of the primary cooling is not restricted as long as the accuracy of the next cooling stop temperature is ensured, but is preferably 700 ° C./second or less in consideration of the current cooling technology. The primary cooling method is not particularly limited, and for example, water cooling by known laminar cooling can be used.
冷却速度30℃/秒以上で行う一次冷却は600〜750℃の温度域で停止させる必要があるが、これは600℃未満の温度ではTiなどの微細炭化物の析出が十分でなく780MPa以上のTSを得るのが困難になり、また、750℃を超える温度ではフェライトや炭化物の粗大化を招き、加工後の伸びフランジ性が低下したり、780MPa以上のTSが得られなくなるためである。なお、冷却停止温度は650〜750℃とすることが好ましい。 The primary cooling performed at a cooling rate of 30 ° C / second or more must be stopped in the temperature range of 600 to 750 ° C, but this is a TS of 780 MPa or more because precipitation of fine carbides such as Ti is not sufficient at temperatures below 600 ° C. This is because it is difficult to obtain a high temperature, and when the temperature exceeds 750 ° C., ferrite and carbides become coarse, the stretch flangeability after processing deteriorates, and a TS of 780 MPa or more cannot be obtained. The cooling stop temperature is preferably 650 to 750 ° C.
一次冷却後の冷却停止時間:0.2〜10秒
一次冷却を行った後は、フェライトの生成量を調整するために冷却を停止して空冷とする必要がある。冷却停止時間が0.2秒未満ではフェライト生成が不十分となり、その結果ベイナイトの生成量が過剰になって加工後の伸びフランジ性が低下し、10秒を超えると炭化物が粗大化し、780MPa以上のTSが得られなくなる。したがって、冷却停止時間(空冷時間)は0.2〜10秒、好ましくは１〜10秒とし、より好ましくは1〜5秒、さらに好ましくは2〜5秒とする。
Cooling stop time after primary cooling: 0.2 to 10 seconds After performing primary cooling, it is necessary to stop cooling to air cooling in order to adjust the amount of ferrite produced. If the cooling stop time is less than 0.2 seconds, ferrite formation is insufficient, resulting in excessive bainite generation and deteriorated stretch flangeability after processing, and if it exceeds 10 seconds, carbides become coarse and TS of 780 MPa or more Cannot be obtained. Therefore, the cooling stop time (air cooling time) is 0.2 to 10 seconds, preferably 1 to 10 seconds, more preferably 1 to 5 seconds, and further preferably 2 to 5 seconds.
冷却停止後の二次冷却条件:冷却速度55℃/秒以上、少なくとも500℃以下の温度域は冷却速度120℃/秒以上で核沸騰冷却となる条件で冷却
冷却停止後は、冷却停止中に調整されたフェライトの生成量が変動しないように、冷却速度55℃/秒以上、好ましくは70℃/秒以上で巻取温度まで二次冷却する必要がある。ここで、二次冷却するにあたっては、水冷を利用するが、本発明においては、500℃以下の温度域で水冷する場合、膜沸騰から核沸騰への遷移沸騰が起こりやすく、温度ムラの問題が生じる。このとき、少なくとも膜沸騰冷却と核沸騰冷却が共存する遷移沸騰冷却となる500℃以下の温度域を冷却速度120℃/秒以上、好ましくは250℃/秒以上で核沸騰冷却となる条件で冷却すれば、温度ムラを確実に解消でき、鋼板内材質変動を安定して小さくすることができる。また、加工後の伸びフランジ性を向上させることができる。なお、核沸騰冷却する場合、前記冷却停止後核沸騰冷却を行うまでの平均冷却速度を55℃/秒以上とし、引き続き核沸騰冷却を開始すればよい。また、上記核沸騰冷却となる条件での冷却における冷却速度は、核沸騰冷却中の平均冷却速度である。本発明においては、少なくとも500℃以下の温度域を核沸騰冷却すればよく、冷却停止後500℃以上の温度から核沸騰冷却となる条件での冷却を開始してもよい。
Secondary cooling conditions after stopping cooling: Cooling at 55 ° C / sec or more, at least 500 ° C or less in conditions where nucleate boiling cooling is performed at a cooling rate of 120 ° C / sec or more. It is necessary to perform secondary cooling to the coiling temperature at a cooling rate of 55 ° C./second or more, preferably 70 ° C./second or more so that the adjusted ferrite production amount does not fluctuate. Here, in the secondary cooling, water cooling is used, but in the present invention, when water cooling is performed at a temperature range of 500 ° C. or lower, transition boiling from film boiling to nucleate boiling is likely to occur, and there is a problem of temperature unevenness. Arise. At this time, at least a temperature range of 500 ° C. or less, which is transition boiling cooling in which film boiling cooling and nucleate boiling cooling coexist, is cooled under a condition of nucleate boiling cooling at a cooling rate of 120 ° C./second or more, preferably 250 ° C./second or more. By doing so, the temperature unevenness can be surely eliminated, and the material variation in the steel plate can be stably reduced. Moreover, the stretch flangeability after a process can be improved. In the case of nucleate boiling cooling, the average cooling rate until the nucleate boiling cooling after the cooling stop is set to 55 ° C./second or more, and then the nucleate boiling cooling is started. Further, the cooling rate in the cooling under the above-described nucleate boiling cooling is the average cooling rate during the nucleate boiling cooling. In the present invention, nucleate boiling cooling may be performed at least in a temperature range of 500 ° C. or lower, and cooling may be started under conditions of nucleate boiling cooling from a temperature of 500 ° C. or higher after cooling is stopped.
特許文献4には、膜沸騰冷却と核沸騰冷却が共存する遷移沸騰冷却される温度域を核沸騰冷却で行い、遷移沸騰冷却によって生じる鋼板面内の温度ムラを縮小させる技術が開示されている。しかし、本発明者らの検討によれば、遷移沸騰冷却を単に核沸騰冷却に代えるだけでは必ずしも鋼板面内の温度ムラを解消できず、安定して温度ムラを解消するには核沸騰冷却時の冷却速度を120℃/秒以上にすることが必要であることを見出した。この理由は、必ずしも明確ではないが、冷却速度が遅いと鋼板表面の水膜の破壊が十分でない部分が発生するためと考えられる。さらに、冷却速度が120℃/秒以上では、加工後の伸びフランジ性の向上が認められる。これは、粒界に微細に析出するセメンタイトの生成を抑制できるためと考えられる。 Patent Document 4 discloses a technique for reducing temperature unevenness in a steel sheet surface caused by transition boiling cooling by performing a transition boiling cooling temperature range in which film boiling cooling and nucleate boiling cooling coexist with nucleate boiling cooling. . However, according to the study by the present inventors, it is not always possible to eliminate the temperature unevenness in the steel sheet surface by simply replacing the transition boiling cooling with the nucleate boiling cooling. It was found that it was necessary to increase the cooling rate of 120 ° C./second or more. The reason for this is not necessarily clear, but it is considered that when the cooling rate is low, a portion where the water film on the steel sheet surface is not sufficiently broken is generated. Furthermore, when the cooling rate is 120 ° C./second or more, improvement in stretch flangeability after processing is recognized. This is considered to be because the production of cementite finely precipitated at the grain boundaries can be suppressed.
冷却速度120℃/秒以上で核沸騰冷却を確実に行うためには、水量密度を2000L/min.m2以上とすることが好ましい。核沸騰冷却を実施するには、従来の方法、すなわち鋼板上面に対しては直進性に優れたラミナーもしくはジェット冷却が好ましい。ノズルの形状としては、一般的に円管やスリットノズルがあるがどちらを採用しても問題はない。また、ラミナーもしくはジェット冷却の流速は4m/秒以上とすることが好ましい。これは、冷却時に鋼板上に生成する液膜をラミナーもしくはジェット冷却により安定的に突き破るための運動量を得る必要があるためである。なお、鋼板下面に対しては重力により冷却水は落下するため、鋼板面に液膜ができないため、スプレー冷却を用いても問題ない。もちろん、鋼板上面の場合と同様なラミナーやジェット冷却を採用することもできる。 In order to reliably perform nucleate boiling cooling at a cooling rate of 120 ° C./sec or more, the water density is preferably 2000 L / min.m 2 or more. In order to carry out nucleate boiling cooling, a conventional method, that is, laminar or jet cooling excellent in linearity with respect to the upper surface of the steel sheet is preferable. As the shape of the nozzle, there are generally a circular tube and a slit nozzle, but there is no problem even if either is adopted. The laminar or jet cooling flow rate is preferably 4 m / second or more. This is because it is necessary to obtain a momentum for stably breaking through the liquid film formed on the steel plate during cooling by laminar or jet cooling. In addition, since cooling water falls with respect to the steel plate lower surface by gravity, since a liquid film cannot be formed on the steel plate surface, there is no problem even if spray cooling is used. Of course, the same laminar and jet cooling as in the case of the upper surface of the steel plate can be employed.
二次冷却後の巻取温度:350〜500℃
二次冷却後まで維持された残留オーステナイトをベイナイトに変態させるために、350〜500℃、好ましくは400〜500℃の巻取温度でコイル状に巻取る必要がある。これは、巻取温度が300℃未満ではベイナイトより硬質なマルテンサイトが生成し、また、500℃を超えるとパーライトが生成して、加工後の伸びフランジ性が低下するためである。巻取温度が500℃未満の場合は、500℃以下では遷移沸騰冷却が起こるため、上述のように、500℃から巻取温度までの温度域を平均冷却速度120℃/秒以上で核沸騰冷却する必要がある。なお、巻取温度が500℃のときは、500℃以下の水冷が不要になるので、核沸騰冷却を考慮する必要はなく、巻取温度までの二次冷却の平均冷却速度を55℃/秒以上とすればよい。
Winding temperature after secondary cooling: 350-500 ° C
In order to transform the retained austenite maintained after the secondary cooling into bainite, it is necessary to wind it in a coil shape at a coiling temperature of 350 to 500 ° C, preferably 400 to 500 ° C. This is because martensite harder than bainite is generated when the coiling temperature is less than 300 ° C., and pearlite is generated when the coiling temperature exceeds 500 ° C., and stretch flangeability after processing is deteriorated. When the coiling temperature is less than 500 ° C, transition boiling cooling occurs at 500 ° C or less. Therefore, as described above, the temperature range from 500 ° C to the coiling temperature is nucleate boiling cooling at an average cooling rate of 120 ° C / second or more. There is a need to. When the coiling temperature is 500 ° C, water cooling below 500 ° C is unnecessary, so there is no need to consider nucleate boiling cooling, and the average cooling rate of secondary cooling up to the coiling temperature is 55 ° C / sec. That is all.
このようにして得られる高強度熱延鋼板の組織は、体積率で80〜95%のフェライトと、ベイナイトあるいはさらに不可避的に生じる3%未満の他の相(マルテンサイトや残留オーステナイト)とから構成されている。 The structure of the high-strength hot-rolled steel sheet thus obtained is composed of 80-95% ferrite by volume and bainite or other inevitably less than 3% other phases (martensite and retained austenite). Has been.
その他の製造条件は通常の条件で行える。例えば、所望の成分組成を有する鋼は転炉や電気炉などで溶製後、真空脱ガス炉にて2次精錬を行って製造される。その後の鋳造は、生産性や品質上の点から連続鋳造法で行うのが好ましい。鋳造後は、本発明の方法にしたがって熱間圧延を行う。熱間圧延後は、表面にスケールが付着した状態であっても、酸洗を行いスケールを除去した状態であっても、鋼板の特性が変わることはない。また、熱間圧延後、調質圧延を行ったり、溶融亜鉛めっき、電気めっき、化成処理を施すことも可能である。 Other manufacturing conditions can be performed under normal conditions. For example, steel having a desired component composition is manufactured by melting in a converter or electric furnace and then performing secondary refining in a vacuum degassing furnace. The subsequent casting is preferably performed by a continuous casting method from the viewpoint of productivity and quality. After casting, hot rolling is performed according to the method of the present invention. After hot rolling, the properties of the steel sheet do not change even if the scale is attached to the surface or the scale is removed by pickling. Further, after hot rolling, temper rolling, hot dip galvanizing, electroplating, and chemical conversion treatment can be performed.
表1に示す化学組成の鋼片を、1250℃に加熱し、表2に示す仕上温度で熱間圧延して板厚3.2mmの熱延板とした後、水冷を利用して表2に示す冷却条件により冷却し、表2に示す巻取温度でコイル状に巻取った。なお、ここで、巻取温度は鋼帯の幅方向中央部の巻取温度を鋼帯の長手方向に計測し、それらを平均した値である。また、このとき、巻取り装置の直前に鋼板表面温度を2次元的に測定可能な放射温度計[NEC三栄(株)製型式TH7800]を設置し、次のように鋼板面の温度ムラを評価した。
温度ムラ:放射温度計で計測された局所的に巻取温度が350℃未満となる低温部の面積を求め、その鋼板の全面積に占めるの割合S[=(低温部の面積)/(鋼板の全面積)×100(%)]を算出し、S<5%であれば温度ムラがないとした。
Steel slabs having the chemical composition shown in Table 1 were heated to 1250 ° C. and hot rolled at a finishing temperature shown in Table 2 to form a hot-rolled sheet having a thickness of 3.2 mm, and then shown in Table 2 using water cooling. The sample was cooled under the cooling conditions and wound into a coil at the winding temperature shown in Table 2. Here, the coiling temperature is a value obtained by measuring the coiling temperature at the center in the width direction of the steel strip in the longitudinal direction of the steel strip and averaging them. At this time, a radiation thermometer [NEC Sanei Model TH7800], which can measure the steel sheet surface temperature in two dimensions, was installed just before the winding device, and the temperature unevenness of the steel sheet surface was evaluated as follows. did.
Temperature non-uniformity: Obtain the area of the low-temperature part where the coiling temperature is locally measured by a radiation thermometer, and the ratio of the total area of the steel sheet S [= (area of the low-temperature part) / (steel sheet Total area) × 100 (%)], and if S <5%, there was no temperature unevenness.
なお、表2中、鋼板No.5およびNo.17は巻取温度が500℃以上であり、核沸騰冷却となる条件での冷却は行わなかった。 In Table 2, the steel plates No. 5 and No. 17 had a coiling temperature of 500 ° C. or higher, and were not cooled under the conditions for nucleate boiling cooling.
次に、巻取り後の熱延板を酸洗後、コイル先端部から長手方向に30m入った位置で鋼板の幅方向中央におけるベイナイト体積率を、また、コイル先端部から30mの位置で幅方向中央、幅方向1/4および3/4の3箇所からJIS 5号引張試験片(圧延方向に直角方向)および穴広げ試験用試験片を採取してTSおよび加工後の穴広げ率λを、次のようにして測定した。
ベイナイト体積率:走査型電子顕微鏡(SEM)用試験片を採取し、圧延方向に平行な板厚断面を研磨後、ナイタール腐食し、倍率1000倍でSEM写真を10視野で撮影し、ベイナイトを画像処理により抽出し、画像解析処理によりベイナイトの面積および観察視野の面積を測定してベイナイト面積率[=(ベイナイトの面積)/(観察視野の面積)×100(%)]を求め、これをベイナイト体積率とした。なお、本発明例では、ベイナイト以外の部分はほぼフェライトであり、フェライト、ベイナイト以外の他の相は3%未満であった。
TS:3本の引張試験片に歪み速度10mm/minで引張試験を行って引張強度TSを求め、3本の平均値をTSとした。
λ:採取した3個の穴広げ試験用試験片に圧下率10%の冷間圧延を施した後、130mm角の板を切り出し、板中央に10mmφの穴を打ち抜いた後、60°円錐ポンチをバリと反対側から押し上げ、亀裂が板厚を貫通した時点での穴径dmmを測定し、次式より算出し、3個の平均値によりλを評価した。
λ(%)=[(d-10)/10]×100
さらに、鋼板内材質変動を調査するため、コイル先端部から長手方向に100、200、400、600、700m入った各位置で、圧延方向に平行な方向を試験片の長手方向として、鋼板の幅方向に、幅方向の両端25mmの内側から25本の試験片を等間隔に採取し、合計125本のJIS 5号引張試験片(圧延方向に平行な方向が引張方向)を採取し、上記と同様な方法でTSを求め、その標準偏差σを算出した。
Next, after pickling the hot-rolled sheet after winding, the bainite volume ratio at the center in the width direction of the steel sheet at a position 30 m in the longitudinal direction from the coil tip, and the width direction at a position 30 m from the coil tip JIS No. 5 tensile test specimens (perpendicular to the rolling direction) and hole expansion test specimens were collected from the center, width direction 1/4 and 3/4, and TS and hole expansion ratio λ after processing were obtained. Measurement was performed as follows.
Bainite volume fraction: A specimen for a scanning electron microscope (SEM) was collected, the thickness cross section parallel to the rolling direction was polished, then it was corroded by Nital, and SEM photographs were taken at 1000 magnifications with 10 fields of view, and bainite was imaged. The area of bainite and the area of the observation visual field are measured by image processing, and the area ratio of bainite [= (area of bainite) / (area of the observation visual field) × 100 (%)] is obtained. The volume ratio was used. In the examples of the present invention, the portion other than bainite was almost ferrite, and the phases other than ferrite and bainite were less than 3%.
TS: A tensile test was performed on the three tensile test pieces at a strain rate of 10 mm / min to obtain the tensile strength TS, and the average value of the three was set as TS.
λ: After performing cold rolling with a reduction rate of 10% on the three test specimens for hole expansion test, a 130 mm square plate was cut out, a 10 mmφ hole was punched in the center of the plate, and then a 60 ° conical punch was Pushed up from the opposite side of the burr, the hole diameter dmm was measured when the crack penetrated the plate thickness, calculated from the following equation, and λ was evaluated by the average value of the three.
λ (%) = [(d-10) / 10] × 100
Furthermore, in order to investigate the material fluctuations in the steel sheet, the width of the steel sheet, with the direction parallel to the rolling direction being the longitudinal direction of the test piece at each position 100, 200, 400, 600, 700 m in the longitudinal direction from the coil tip. In the direction, 25 test pieces from the inner side of 25 mm in the width direction are collected at regular intervals, and a total of 125 JIS No. 5 tensile test pieces (the direction parallel to the rolling direction is the tensile direction) are collected. TS was calculated | required by the same method and the standard deviation (sigma) was computed.
結果を表3に示す。本発明例では、TSが780MPa以上であり、かつ加工後平均λが80%以上で加工後の伸びフランジ性にも優れているとともに、コイル内の温度ムラがほとんどないためTSの標準偏差σは15MPa以下と小さく、鋼板内材質変動が小さいことがわかる。 The results are shown in Table 3. In the example of the present invention, TS is 780 MPa or more, the average λ after processing is 80% or more and excellent in stretch flangeability after processing, and there is almost no temperature unevenness in the coil, so the standard deviation σ of TS is It can be seen that the material fluctuation in the steel sheet is small, as small as 15 MPa or less.
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| JP2009191360A (en) * | 2008-01-17 | 2009-08-27 | Jfe Steel Corp | High strength steel sheet, and method for producing the same |
| JP2011068945A (en) * | 2009-09-25 | 2011-04-07 | Jfe Steel Corp | High strength hot rolled steel sheet, and method for producing the same |
| JP2011122189A (en) * | 2009-12-09 | 2011-06-23 | Jfe Steel Corp | HIGH STRENGTH HOT ROLLED STEEL SHEET EXCELLENT IN ELONGATION AND STRETCH-FLANGE PROPERTY AND HAVING TENSILE STRENGTH OF >=780 MPa, AND METHOD FOR PRODUCING IT |
| JP2011140699A (en) * | 2010-01-08 | 2011-07-21 | Jfe Steel Corp | Method for manufacturing high-strength hot-rolled steel sheet excellent in elongation property and stretch flanging property |
| KR101148921B1 (en) | 2010-01-28 | 2012-05-22 | 현대제철 주식회사 | Cold rolled steel sheet and manufacturing method thereof |
| WO2016143298A1 (en) * | 2015-03-06 | 2016-09-15 | Ｊｆｅスチール株式会社 | High strength steel sheet and manufacturing method therefor |
| CN113106332A (en) * | 2021-02-24 | 2021-07-13 | 唐山不锈钢有限责任公司 | Low-crack-sensitivity 510L hot-rolled automobile structure steel strip and production method thereof |
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| JP2009191360A (en) * | 2008-01-17 | 2009-08-27 | Jfe Steel Corp | High strength steel sheet, and method for producing the same |
| JP2011068945A (en) * | 2009-09-25 | 2011-04-07 | Jfe Steel Corp | High strength hot rolled steel sheet, and method for producing the same |
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| JP2011140699A (en) * | 2010-01-08 | 2011-07-21 | Jfe Steel Corp | Method for manufacturing high-strength hot-rolled steel sheet excellent in elongation property and stretch flanging property |
| KR101148921B1 (en) | 2010-01-28 | 2012-05-22 | 현대제철 주식회사 | Cold rolled steel sheet and manufacturing method thereof |
| WO2016143298A1 (en) * | 2015-03-06 | 2016-09-15 | Ｊｆｅスチール株式会社 | High strength steel sheet and manufacturing method therefor |
| JPWO2016143298A1 (en) * | 2015-03-06 | 2017-04-27 | Ｊｆｅスチール株式会社 | High strength steel plate and manufacturing method thereof |
| US10815547B2 (en) | 2015-03-06 | 2020-10-27 | Jfe Steel Corporation | High strength steel sheet and manufacturing method therefor |
| CN113106332A (en) * | 2021-02-24 | 2021-07-13 | 唐山不锈钢有限责任公司 | Low-crack-sensitivity 510L hot-rolled automobile structure steel strip and production method thereof |
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