JP5482205B2 - High strength hot rolled steel sheet and method for producing the same - Google Patents
High strength hot rolled steel sheet and method for producing the same Download PDFInfo
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- 229910000831 Steel Inorganic materials 0.000 title claims description 48
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- 238000005098 hot rolling Methods 0.000 claims description 21
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- 229910052748 manganese Inorganic materials 0.000 claims description 5
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- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 6
- 229910052725 zinc Inorganic materials 0.000 description 6
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Description
本発明は、自動車の構造部材などに有用な高強度熱延鋼板、特に、引張強さTSが690〜980MPa、全伸びElが27%以上、伸びフランジ性の指標となる穴拡げ率λが50%以上で、かつ鋼板内におけるTSの均一性に優れた高強度熱延鋼板およびその製造方法に関する。 The present invention is a high-strength hot-rolled steel sheet useful for structural members of automobiles, in particular, a tensile strength TS of 690 to 980 MPa, a total elongation El of 27% or more, and a hole expansion ratio λ that is an index of stretch flangeability is 50. The present invention relates to a high-strength hot-rolled steel sheet that is at least% and excellent in TS uniformity in the steel sheet and a method for producing the same.
近年、地球環境保全の観点からCO2の排出量の規制が厳しくなるなか、自動車の燃費改善が急務とされ、使用される部材の薄肉化による軽量化が要求されている。加えて、自動車の衝突時に乗員の安全を確保するため、使用される部材の高強度化も要求されている。このため、自動車車体の軽量化および高強度化の双方が積極的に進められている。 In recent years, as regulations on CO 2 emissions have become stricter from the viewpoint of global environmental conservation, there is an urgent need to improve the fuel efficiency of automobiles, and weight reduction is required by reducing the thickness of the components used. In addition, in order to ensure the safety of passengers in the event of a car collision, it is also required to increase the strength of the members used. For this reason, both reduction in weight and strength of the automobile body are being actively promoted.
現在では、乗用車のピラーやメンバー、トラックのフレームなどの構造部材に主として440MPa級や590MPa級のTSを有する高強度熱延鋼板が使用されるようになっているが、近い将来、690〜980MPaのTSを有する高強度熱延鋼板の実用化が予測されている。そのため、こうした強度レベルの高強度熱延鋼板を対象とした技術開発が活発に行われており、高強度化にともなって劣化する加工性の向上、なかでも伸び特性や伸びフランジ性などの延性の向上を図った種々の高強度熱延鋼板が提案されている。例えば、特許文献1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%以下およびTi:0.03〜0.20%を含有し、残部がFeおよび不可避的不純物よりなる成分組成を有するとともに、体積占有率で50〜90%がフェライト相で、かつ残部が実質的にベイナイト相であって、フェライト相とベイナイト相の体積占有率の合計が95%以上であり、フェライト相中には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TSが780MPa以上の高強度熱延鋼板が開示されている。また、特許文献2には、質量%で、C:0.015〜0.06%、Si:0.5%未満、Mn:0.1〜2.5%、P:0.10%以下、S:0.01%以下、Al:0.005〜0.3%、N:0.01%以下、Ti:0.01〜0.30%、B:2〜50ppmを含有し、残部Feおよび不可避的不純物からなる鋼組成を有し、0.75<(C/12)/(Ti/48-N/14-S/32)<1.25および1.0<Mn+B/10-Siの関係を満足し、フェライト相とベイニティックフェライト相の面積率の合計が90%以上、セメンタイトの面積率が5%以下であり、TSが690〜850MPa、λが40%以上である伸びフランジ性に優れた高強度熱延鋼板が開示されている。さらに、特許文献3には、質量%で、C:0.1%以下、Mo:0.05〜0.6%、Ti:0.02〜0.10%を含み、実質的にフェライト組織に原子比でTi/Mo≧0.1を満たす範囲でTiおよびMoを含む炭化物が分散析出してなる材質均一性に優れたTSが610〜830MPaの高成形性高張力熱延鋼板が開示されている。
Currently, high-strength hot-rolled steel sheets with 440MPa class and 590MPa class TS are mainly used for structural members such as pillars and members of passenger cars and truck frames. The practical application of high-strength hot-rolled steel sheets with TS is expected. Therefore, technological development for high-strength hot-rolled steel sheets with such strength levels has been actively carried out, improving workability that deteriorates with increasing strength, especially ductility such as stretch characteristics and stretch flangeability. Various high-strength hot-rolled steel sheets that have been improved have been proposed. For example, Patent Document 1 includes 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, and Ti: 0.03 to 0.20% contained, the balance has a component composition consisting of Fe and inevitable impurities, 50 to 90% by volume occupancy is the ferrite phase, and the balance is substantially the bainite phase, The total volume occupancy of the ferrite phase and bainite phase is 95% or more, precipitates containing Ti are precipitated in the ferrite phase, and the average diameter of the precipitates is 20 nm or less, and the steel A high-strength hot-rolled steel sheet having a TS of 780 MPa or more, which is excellent in elongation characteristics, stretch flange characteristics, and tensile fatigue characteristics, in which 80% or more of the Ti content is precipitated, is disclosed. Further, in
しかしながら、特許文献1に記載の高強度熱延鋼板では、同文献に記載された製造方法で製造すると鋼板内において均一なTSが安定して得られない場合があるという問題がある。特許文献2や3に記載の高強度熱延鋼板では、Elが低く、また必ずしも高いλが得られないという問題がある。
However, the high-strength hot-rolled steel sheet described in Patent Document 1 has a problem that uniform TS may not be stably obtained in the steel sheet when manufactured by the manufacturing method described in the same document. The high-strength hot-rolled steel sheets described in
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、TSが690〜980MPa、Elが27%以上、λが50%以上で、かつ鋼板内におけるTSのばらつきΔTSが安定して15MPa以下となる高強度熱延鋼板およびそ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. TS is 690 to 980 MPa, El is 27% or more, λ is 50% or more, and TS variation ΔTS in the steel sheet is stably 15 MPa. An object of the present invention is to provide a high-strength hot-rolled steel sheet and a method for producing the same.
本発明者らは、上記の目的とする高強度熱延鋼板について検討を重ねた結果、以下のことを見出した。 As a result of repeated studies on the above-described high strength hot-rolled steel sheet, the present inventors have found the following.
i) 成分組成の適正化を図った上で、フェライト相と、マルテンサイト相を主体とした第二相とからなり、組織全体に占めるフェライト相の面積率が65〜80%で、フェライト相とマルテンサイト相の合計の面積率が95%以上であり、フェライト相と第二相のビッカース硬度差ΔHvが250以下であるミクロ組織にすることにより、690〜980MPaのTSで、27%以上のElと50%以上のλが得られる。 i) After optimizing the component composition, it consists of a ferrite phase and a second phase mainly composed of a martensite phase. The area ratio of the ferrite phase occupying the entire structure is 65 to 80%. By making the microstructure in which the total area ratio of the martensite phase is 95% or more and the Vickers hardness difference ΔHv between the ferrite phase and the second phase is 250 or less, at a TS of 690 to 980 MPa, El of 27% or more Λ of 50% or more is obtained.
ii) 鋼板内において均一なTSを確保してTSのばらつきΔTSを小さくするためには、フェライト相の面積率のばらつきΔSFを小さくすることが重要である。ΔSFを小さくする上では、フェライトフォーマーであるSiの含有量を比較的低め、具体的にはSi≦0.1質量%とした上で、熱間圧延時の仕上温度を比較的高い温度、具体的には(Ar3変態点+100℃)以上として熱間圧延を終了することが重要である。本発明では、熱間圧延後の冷却に引き続いて行う空冷時にフェライト変態させるが、仕上温度を高温とすることにより、フェライト変態の空冷時間感受性を鈍化させ、フェライト相の面積率を確保しつつ、ΔSFを2%以下として、TSのばらつきΔTSを15MPa以下にし、均一なTSを確保する。なお、本発明において、ΔSF、ΔTSは、それぞれ後述する方法にて測定した鋼板内のフェライト相の面積率、TSの測定値の標準偏差(σ)である。 ii) In order to secure uniform TS in the steel sheet and reduce the TS variation ΔTS, it is important to reduce the ferrite phase area ratio variation ΔSF. In order to reduce ΔSF, the content of Si, which is a ferrite former, is relatively low, specifically, Si ≦ 0.1% by mass, and the finishing temperature during hot rolling is relatively high. For this, it is important to finish the hot rolling at (Ar 3 transformation point + 100 ° C.) or more. In the present invention, the ferrite transformation is performed during air cooling subsequent to cooling after hot rolling, but by making the finishing temperature high, the air cooling time sensitivity of the ferrite transformation is dulled, while ensuring the area ratio of the ferrite phase, ΔSF is set to 2% or less, TS variation ΔTS is set to 15 MPa or less, and uniform TS is ensured. In the present invention, ΔSF and ΔTS are the area ratio of the ferrite phase in the steel sheet and the standard deviation (σ) of the measured value of TS, respectively, measured by the method described later.
本発明は、このような知見に基づいてなされたもので、質量%で、C:0.060〜0.150%、Si:0.1%以下、Mn:0.8〜1.8%、P:0.030%以下、S:0.005%以下、Al:0.005〜0.1%、N:0.005%以下、Ti:0.032〜0.120%を含み、残部がFeおよび不可避的不純物からなり、下記の式(1)および式(2)を満足する成分組成を有し、フェライト相と、マルテンサイト相を含む第二相とからなり、組織全体に占める前記フェライト相の面積率が65〜80%で、組織全体に占める前記フェライト相と前記マルテンサイト相の合計の面積率が95%以上であり、前記フェライト相の面積率のばらつきΔSFが2％以下で、前記フェライト相と前記第二相のビッカース硬度差ΔHvが250以下であるミクロ組織を有することを特徴とする高強度熱延鋼板を提供する。
0.05≦C*≦0.09・・・(1)
0.03≦Ti*≦0.10・・・(2)
ただし、C*=[C]-0.55×Ti*、Ti*=[Ti]-48×[N]/14、[M]は元素Mの含有量(質量%)を表す。
The present invention was made based on such findings, and in mass%, C: 0.060 to 0.150%, Si: 0.1% or less, Mn: 0.8 to 1.8%, P: 0.030% or less, S: 0.005% Hereinafter, Al: 0.005 to 0.1%, N: 0.005% or less, Ti: 0.032 to 0.120%, the balance is composed of Fe and unavoidable impurities, and satisfies the following formula (1) and formula (2) The ferrite phase and the second phase containing the martensite phase, the area ratio of the ferrite phase in the entire structure is 65 to 80%, the ferrite phase and the martensite phase in the entire structure The total area ratio is 95% or more, the variation in area ratio ΔSF of the ferrite phase is 2% or less, and the microstructure has a Vickers hardness difference ΔHv between the ferrite phase and the second phase of 250 or less. A high-strength hot-rolled steel sheet is provided.
0.05 ≦ C * ≦ 0.09 ... (1)
0.03 ≦ Ti * ≦ 0.10 ... (2)
However, C * = [C] −0.55 × Ti * , Ti * = [Ti] −48 × [N] / 14, [M] represents the content (mass%) of the element M.
本発明の高強度熱延鋼板は、上記の成分組成を有する鋼スラブを、1200〜1300℃の加熱温度で加熱し、(Ar3変態点+100℃)以上の仕上温度で熱間圧延後、2秒以内に25℃/秒以上の平均冷却速度で600〜720℃の冷却停止温度まで一次冷却し、引き続き5〜60秒間空冷後、25℃/秒以上の平均冷却速度で二次冷却し、400℃未満の巻取温度で巻取ることによって製造できる。 The high-strength hot-rolled steel sheet of the present invention is a steel slab having the above component composition, heated at a heating temperature of 1200 to 1300 ° C, after hot rolling at a finishing temperature of (Ar 3 transformation point + 100 ° C) or higher, Within 2 seconds, primary cooling is performed at an average cooling rate of 25 ° C / second or more to a cooling stop temperature of 600 to 720 ° C, followed by air cooling for 5 to 60 seconds, followed by secondary cooling at an average cooling rate of 25 ° C / second or more, It can be manufactured by winding at a winding temperature of less than 400 ° C.
本発明により、TSが690〜980MPa、Elが27%以上、λが50%以上で、かつ鋼板内におけるΔTSが15MPa以下となる高強度熱延鋼板が製造可能になった。本発明の高強度熱延鋼板は、乗用車のピラーやメンバー、トラックのフレームなどの構造部材に好適である。 According to the present invention, a high-strength hot-rolled steel sheet having TS of 690 to 980 MPa, El of 27% or more, λ of 50% or more, and ΔTS in the steel sheet of 15 MPa or less can be produced. The high-strength hot-rolled steel sheet of the present invention is suitable for structural members such as pillars and members of passenger cars and truck frames.
以下に、本発明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。 Details of the present invention will be described below. Note that “%” representing the content of each component element means “% by mass” unless otherwise specified.
1) 成分組成
C:0.060〜0.150%
Cは、マルテンサイト相の生成を促進したり、フェライト相中に微細なTiの炭化物を形成して高強度化に寄与する元素である。690MPa以上のTSを得るためにはC量を0.060％以上とする必要がある。一方、C量が0.150％を超えるとElやλが低下するのみならず、靱性も劣化する。したがって、C量は0.060〜0.150％とする。
1) Component composition
C: 0.060-0.150%
C is an element that promotes the formation of a martensite phase or contributes to an increase in strength by forming fine Ti carbides in the ferrite phase. In order to obtain a TS of 690 MPa or more, the C content needs to be 0.060% or more. On the other hand, when the amount of C exceeds 0.150%, not only El and λ are lowered, but also toughness is deteriorated. Therefore, the C content is 0.060 to 0.150%.
なお、本発明において高強度化に寄与するマルテンサイト相の量(面積率)を確保するには、Tiの炭化物を形成する以外のC量、すなわちC*=[C]-0.55×Ti*を0.05%以上にする必要がある。ここで、Ti*は[Ti]-48×[N]/14と定義され、炭化物として析出できるTi量を表している。また、[C]-0.55×Ti*としたのは、TiCやTi4C2S2などのTi系析出物として消費されるC量を除いて第二相に濃化する有効C量を考慮するためである。一方、C*が0.09%を超えるとマルテンサイト相の硬度が高くなり過ぎ、フェライト相と第二相のビッカース硬度差を小さくできなくなる。したがって、上記の式(1)のように、0.05≦C*≦0.09とする。 In order to secure the amount (area ratio) of the martensite phase that contributes to high strength in the present invention, the amount of C other than that of forming Ti carbide, that is, C * = [C] −0.55 × Ti * . It must be 0.05% or more. Here, Ti * is defined as [Ti] −48 × [N] / 14, and represents the amount of Ti that can be precipitated as carbide. In addition, [C] -0.55 × Ti * is taken into account the amount of effective C concentrated in the second phase excluding the amount of C consumed as Ti-based precipitates such as TiC and Ti 4 C 2 S 2 It is to do. On the other hand, if C * exceeds 0.09%, the hardness of the martensite phase becomes too high, and the Vickers hardness difference between the ferrite phase and the second phase cannot be reduced. Therefore, 0.05 ≦ C * ≦ 0.09, as in the above formula (1).
Si:0.1%以下
Si量が0.1%を超えるとAr3変態点が上昇し過ぎるため、熱間圧延時に仕上温度、すなわち(Ar3変態点+100℃)以上の確保が困難になり、鋼板内においてTSの不均一を引き起こす。また、Si量が増加すると靱性や耐疲労特性の劣化にもつながる。したがって、Si量は0.1%以下、好ましくは0.05%以下とする。
Si: 0.1% or less
If the Si content exceeds 0.1%, the Ar 3 transformation point will rise too much, making it difficult to ensure a finish temperature of hot rolling, that is, (Ar 3 transformation point + 100 ° C) or higher, and non-uniformity of TS in the steel sheet. cause. In addition, increasing the amount of Si leads to deterioration of toughness and fatigue resistance. Therefore, the Si content is 0.1% or less, preferably 0.05% or less.
Mn:0.8〜1.8%
Mnは、高強度化に有効であるとともに、Ar3変態点を下げ、フェライト粒を微細化させて伸び特性などの延性を向上させる効果を有する。こうした効果を得るにはMn量を0.8%以上とする必要がある。一方、Mn量が1.8%を超えると延性が低下したり、Tiの炭化物の析出が不安定になって鋼板内におけるTSの不均一を増長させる。したがって、Mn量は0.8〜1.8%、好ましくは1.0〜1.8%、より好ましくは1.2〜1.8%とする。
Mn: 0.8-1.8%
Mn is effective for increasing the strength and has the effect of lowering the Ar 3 transformation point and making ferrite grains finer to improve ductility such as elongation characteristics. In order to obtain such an effect, the Mn content needs to be 0.8% or more. On the other hand, if the Mn content exceeds 1.8%, the ductility is lowered, or the precipitation of Ti carbide becomes unstable, and the nonuniformity of TS in the steel sheet is increased. Therefore, the amount of Mn is 0.8 to 1.8%, preferably 1.0 to 1.8%, more preferably 1.2 to 1.8%.
P:0.030%以下
P量が0.030%を超えると粒界への偏析が顕著になり、靱性や溶接性の劣化を招く。したがって、P量は0.030%以下とするが、極力低減することが好ましい。
P: 0.030% or less
When the P content exceeds 0.030%, segregation to the grain boundary becomes remarkable, leading to deterioration of toughness and weldability. Therefore, the P content is 0.030% or less, but it is preferable to reduce it as much as possible.
S:0.005%以下
Sは、MnやTiと硫化物を形成し、伸びフランジ性を低下させる。したがって、S量は0.005%以下とするが、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 and reduces stretch flangeability. Therefore, the S amount is 0.005% or less, but it is preferable to reduce it as much as possible.
Al:0.005〜0.1%
Alは、鋼の脱酸剤として添加され、その清浄度を向上させるのに有効な元素である。こうした効果を得るにはAl量を0.005%以上にする必要がある。一方、Al量が0.1%を超えると表面欠陥が生じやすくなるとともに、コスト増を招く。したがって、Al量は0.005〜0.1%とする。
Al: 0.005-0.1%
Al is an element that is added as a deoxidizer for steel and is effective in improving its cleanliness. In order to obtain such effects, the Al content needs to be 0.005% or more. On the other hand, if the Al content exceeds 0.1%, surface defects are likely to occur, and the cost is increased. Therefore, the Al content is 0.005 to 0.1%.
N:0.005%以下
Nは、Tiとの親和力が強い元素である。そのため、N量が0.005%を超えると炭化物を形成し高強度化に寄与するTi量を確保するために多量のTiを添加する必要があり、コスト増を招く。したがって、N量は0.005%以下とするが、極力低減することが好ましい。
N: 0.005% or less
N is an element having a strong affinity for Ti. For this reason, if the N content exceeds 0.005%, it is necessary to add a large amount of Ti in order to ensure the Ti content that forms carbides and contributes to high strength, resulting in an increase in cost. Therefore, the N content is 0.005% or less, but it is preferable to reduce it as much as possible.
Ti:0.032〜0.120%
Tiは、本発明における重要な元素であり、熱間圧延後の一次冷却に続く空冷時にフェライト相中に粒径が20nm未満の微細なTiCやTi4C2S2などの炭化物として析出し、高強度化に寄与する。また、これらの炭化物によりフェライト相が高強度化されるため、フェライト相と第二相のビッカース硬度差を小さくできる。こうした効果を得るにはTi量を0.032%以上とする必要がある。一方、Ti量が0.120%を超えるとスラブ凝固時に析出するTiの炭化物が粗大化するため、熱間圧延に先立つスラブ加熱時にTiの炭化物を溶解しにくくなり、熱間圧延後に高強度化に寄与する微細なTiの炭化物の析出量を確保できなくなったり、炭化物の不均一な析出を引き起こし、鋼板内におけるTSの均一化を阻害する。したがって、Ti量は0.032〜0.120%とする。
Ti: 0.032-0.120%
Ti is an important element in the present invention, and precipitates as fine TiC and Ti 4 C 2 S 2 carbides having a particle size of less than 20 nm in the ferrite phase during air cooling following primary cooling after hot rolling, Contributes to high strength. Moreover, since the strength of the ferrite phase is increased by these carbides, the difference in Vickers hardness between the ferrite phase and the second phase can be reduced. In order to obtain these effects, the Ti content needs to be 0.032% or more. On the other hand, if the Ti content exceeds 0.120%, the Ti carbide that precipitates during slab solidification will coarsen, making it difficult to dissolve the Ti carbide during slab heating prior to hot rolling, contributing to higher strength after hot rolling. As a result, it becomes impossible to secure the amount of fine Ti carbide precipitates, or cause non-uniform precipitation of carbides, which hinders uniformization of TS in the steel sheet. Therefore, the Ti amount is 0.032 to 0.120%.
なお、Tiは炭化物のみならず、高強度化に寄与しない粗大な窒化物としても析出する。そのため、上記のような粒径が20nm未満の微細な炭化物による高強度化を効果的に図るには、炭化物として析出できるTi量であるTi*を0.03%以上にする必要がある。一方、Ti*が0.10%を超えるとスラブ凝固時に粗大なTiの炭化物が析出し、上記のようなスラブ加熱後にも粗大なTi炭化物が残存するという問題を引き起こす。したがって、上記の式(2)のように、0.03≦Ti*≦0.10とする必要がある。 Ti precipitates not only as carbides but also as coarse nitrides that do not contribute to high strength. Therefore, in order to effectively increase the strength with fine carbides having a particle size of less than 20 nm as described above, it is necessary to set Ti * , which is the amount of Ti that can be precipitated as carbides, to 0.03% or more. On the other hand, if Ti * exceeds 0.10%, coarse Ti carbide precipitates during slab solidification, which causes the problem that coarse Ti carbide remains even after slab heating as described above. Therefore, it is necessary to satisfy 0.03 ≦ Ti * ≦ 0.10 as in the above formula (2).
残部はFeおよび不可避的不純物である。 The balance is Fe and inevitable impurities.
2) ミクロ組織
690〜980MPaのTSと優れた伸び特性や伸びフランジ性を両立させるには、フェライト相と、マルテンサイト相を主体とする第二相とからなるミクロ組織にすることが効果的である。これは、延性に富んだフェライト相中にマルテンサイト相のような硬質相を混在させて高強度化を図るとともに、フェライト相と第二相との硬度差を小さくして両相の界面における応力集中を極力緩和し、伸び特性や伸びフランジ性を向上させようという技術思想に基づいている。
2) Micro structure
In order to achieve both 690 to 980 MPa TS and excellent elongation characteristics and stretch flangeability, it is effective to form a microstructure composed of a ferrite phase and a second phase mainly composed of a martensite phase. This is because the ferrite phase rich in ductility is mixed with a hard phase such as a martensite phase to increase the strength, and the hardness difference between the ferrite phase and the second phase is reduced to reduce the stress at the interface between both phases. It is based on the technical idea of reducing concentration as much as possible and improving stretch characteristics and stretch flangeability.
しかし、フェライト相の組織全体に占める面積率が65%未満ではElが27%以上の優れた伸び特性が得られず、80%を超えると第二相にCが濃化して第二相を過度に硬化させるため、フェライト相と第二相との硬度差が大きくなり優れた伸び特性や伸びフランジ性が得られない。また、フェライト相とマルテンサイト相の組織全体に占める合計の面積率が95%未満ではパーライト相などが混在することになり、強度や伸び特性や伸びフランジ特性の安定性の劣化を招く。したがって、フェライト相と、マルテンサイト相を含む第二相とからなり、組織全体に占めるフェライト相の面積率が65〜80%、組織全体に占めるフェライト相とマルテンサイト相の合計の面積率が95%以上であるミクロ組織にする必要がある。 However, if the area ratio of the ferrite phase in the entire structure is less than 65%, excellent elongation characteristics with El of 27% or more cannot be obtained, and if it exceeds 80%, C concentrates in the second phase and the second phase becomes excessive. Therefore, the hardness difference between the ferrite phase and the second phase becomes large, and excellent elongation characteristics and stretch flangeability cannot be obtained. Further, if the total area ratio of the ferrite phase and martensite phase in the entire structure is less than 95%, a pearlite phase or the like is mixed, which leads to deterioration of strength, elongation characteristics, and stability of stretch flange characteristics. Therefore, it consists of a ferrite phase and a second phase containing a martensite phase. The area ratio of the ferrite phase in the entire structure is 65 to 80%, and the total area ratio of the ferrite phase and the martensite phase in the entire structure is 95%. It is necessary to make the microstructure more than%.
さらに、λが50%以上の優れた伸びフランジ性を得るにはフェライト相と第二相のビッカース硬度差ΔHvを250以下にする必要がある。これは、上述したように、フェライト相中に粒径が20nm未満の微細なTiCやTi4C2S2などの炭化物を析出させることにより達成される。 Further, in order to obtain excellent stretch flangeability with λ of 50% or more, the Vickers hardness difference ΔHv between the ferrite phase and the second phase needs to be 250 or less. As described above, this is achieved by precipitating fine carbides such as TiC or Ti 4 C 2 S 2 having a particle size of less than 20 nm in the ferrite phase.
また、フェライト相の面積率のばらつきΔSFは、鋼板内におけるTSの均一性に大きな影響を及ぼす。TSのばらつきΔTSを15MPa以下にしてTSの均一性を確保するには、ΔSFを2%以下にする必要がある。 Further, the variation ΔSF in the area ratio of the ferrite phase greatly affects the uniformity of TS in the steel sheet. In order to ensure TS uniformity with TS variation ΔTS of 15 MPa or less, ΔSF needs to be 2% or less.
なお、第二相には、上記の面積率の条件を満足する範囲内であれば、マルテンサイト相以外にベイナイト相やパーライト相を含んでも本発明の効果が損なわれることはない。 In addition, the effect of this invention is not impaired even if a 2nd phase contains the bainite phase and a pearlite phase other than a martensite phase, if it exists in the range which satisfies said area ratio conditions.
ここで、組織全体に占めるフェライト相やマルテンサイト相の面積率は、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、板厚中心部近傍で倍率1000倍のSEM写真を10視野撮影し、フェライト相やマルテンサイト相を画像処理により識別し、画像解析処理によりフェライト相やマルテンサイト相の面積を測定し、観察視野の面積に占める割合(百分率)として求めた。 Here, the area ratio of the ferrite phase and martensite phase occupying the entire structure is obtained by taking a specimen for a scanning electron microscope (SEM), polishing the plate thickness section parallel to the rolling direction, then corroding the nital, Take 10 SEM photographs at 1000x magnification near the center, identify the ferrite phase and martensite phase by image processing, measure the area of the ferrite phase and martensite phase by image analysis processing, and make the area of the observation field of view It calculated | required as a ratio (percentage).
フェライト相の面積率のばらつきΔSFは、コイル状の鋼板の最内周と最外周の一巻き目をカットし、コイルの幅方向の両端部10mmをトリミング後、コイル長手方向に20等分、幅方向に8等分に分割し、端部も含めた189点の位置から試料を採取し、上記の方法で測定したフェライト相の占有率の標準偏差として求めた。なお、本発明の鋼板のフェライト相の組織全体に占める面積率およびフェライト相とマルテンサイト相の組織全体に占める面積率の合計は、上記189点で観察した面積率の平均値とした。 The variation in the ferrite phase area ratio ΔSF is cut into the innermost and outermost windings of the coiled steel sheet, trimming both ends 10mm in the width direction of the coil, and then dividing the width into 20 equal parts in the coil longitudinal direction. The sample was divided into eight equal parts, and samples were taken from 189 positions including the ends, and obtained as the standard deviation of the ferrite phase occupancy measured by the above method. The total area ratio of the ferrite phase and the total area ratio of the ferrite phase and the martensite phase in the steel sheet of the present invention was the average value of the area ratios observed at 189 points.
フェライト相と第二相の硬度測定は、コイル長手方向および幅方向の中心部より試料を採取し、圧延方向に平行な板厚断面を研磨後ナイタール腐食し、現出した各相について、JIS Z2244(2009)に準じて、測定条件は板厚1/4部近傍にて試験力29.4mN(荷重3g)でフェライト相および第二相をそれぞれ15箇所で測定を行い平均値を各相の硬度とした。このときの平均硬度をフェライト相についてはHvF、第二相についてはHv’として表し、ビッカース硬度差ΔHvを下式にて求めた。 For the hardness measurement of the ferrite phase and the second phase, a sample was taken from the central part in the coil longitudinal direction and width direction, the plate thickness section parallel to the rolling direction was polished and then subjected to Nital corrosion, and for each phase that appeared, JIS Z2244 According to (2009), the measurement conditions were measured at 15 locations each for the ferrite phase and the second phase with a test force of 29.4 mN (load 3 g) in the vicinity of 1/4 part of the plate thickness, and the average value was determined as the hardness of each phase. did. The average hardness at this time was expressed as Hv F for the ferrite phase and Hv ′ for the second phase, and the Vickers hardness difference ΔHv was determined by the following equation.
ΔHv=Hv’-HvF
3) 製造条件
スラブの加熱温度:1200〜1300℃
熱間圧延後フェライト相中に微細なTiの炭化物を析出させるには熱間圧延前にスラブ中に析出している粗大なTiの炭化物を溶解させる必要がある。それにはスラブを1200℃以上に加熱する必要がある。一方、スラブを1300℃を超えて加熱するとスケールの生成が増大し、歩留まりの低下を招く。したがって、スラブの加熱温度は1200〜1300℃とする。
ΔHv = Hv'-Hv F
3) Manufacturing conditions Slab heating temperature: 1200-1300 ℃
In order to precipitate fine Ti carbide in the ferrite phase after hot rolling, it is necessary to dissolve coarse Ti carbide precipitated in the slab before hot rolling. This requires heating the slab to 1200 ° C or higher. On the other hand, when the slab is heated to over 1300 ° C., scale generation increases and yield decreases. Therefore, the heating temperature of the slab is set to 1200 to 1300 ° C.
熱間圧延の仕上温度:(Ar3変態点+100℃)以上
熱間圧延の仕上温度によって圧延直後のオーステナイト粒径や蓄積するひずみエネルギーが変化し、フェライト変態挙動やTiの炭化物の析出挙動に影響を及ぼす。フェライト相の面積率およびそのばらつきを上記のように制御するとともに、フェライト相中に微細なTiの炭化物を析出させるには、フェライト変態を急激に進行させないようにすることが重要であるが、それには圧延直後のオーステナイト相に蓄積されたひずみエネルギーが小さく、オーステナイ相の粒径を比較的大きな状態にする必要がある。このためには、仕上温度をAr3変態点に対して高温とする必要があり、具体的には(Ar3変態点+100℃)以上とする必要がある。なお、ここでAr3変態点は、変態点測定装置を用い、例えば1200℃に10min加熱した後、1℃/秒の冷却速度で冷却する条件により求めればよい。仕上温度が(Ar3変態点+100℃)未満だとひずみエネルギーが蓄積した未再結晶オーステナイト相や微細なオーステナイト粒が増加し、フェライト変態やTiの炭化物の析出が急激に進行してフェライト相の面積率およびそのばらつきやTiの炭化物の析出を制御することが困難になり、伸び特性や伸びフランジ性の低下や鋼板内におけるTSの不均一を引き起こす。したがって、仕上温度は(Ar3変態点+100℃)以上とする。
Hot rolling finishing temperature: (Ar 3 transformation point + 100 ° C) or more The hot rolling finishing temperature changes the austenite grain size immediately after rolling and the accumulated strain energy, resulting in ferrite transformation behavior and Ti carbide precipitation behavior. affect. In order to control the area ratio of the ferrite phase and its variation as described above and to precipitate fine Ti carbide in the ferrite phase, it is important to prevent the ferrite transformation from proceeding rapidly. The strain energy accumulated in the austenite phase immediately after rolling is small, and the grain size of the austenite phase needs to be relatively large. For this purpose, the finish must be a high temperature the temperature with respect to Ar 3 transformation point, and specifically there is a need to (Ar 3 transformation point + 100 ° C.) or higher. Here, the Ar 3 transformation point may be determined by using a transformation point measuring device, for example, by heating at 1200 ° C. for 10 minutes and then cooling at a cooling rate of 1 ° C./second. When the finishing temperature is less than (Ar 3 transformation point + 100 ° C), unrecrystallized austenite phase and fine austenite grains with accumulated strain energy increase, and ferrite transformation and precipitation of Ti carbide progress rapidly and the ferrite phase This makes it difficult to control the area ratio and variation of Ti and precipitation of Ti carbide, causing deterioration of stretch characteristics and stretch flangeability, and non-uniformity of TS in the steel sheet. Therefore, the finishing temperature is (Ar 3 transformation point + 100 ° C.) or higher.
なお、仕上温度を(Ar3変態点+100℃)以上にする上で、Ar3変態点が高いと仕上温度の確保が困難になり、鋼板内におけるTSの不均一の原因となるので、本発明では、上述したように、Si量を低下させてAr3変態点を下げ、仕上温度を確実に確保できるようにしている。 Note that when the finishing temperature is set to (Ar 3 transformation point + 100 ° C) or higher, if the Ar 3 transformation point is high, it becomes difficult to secure the finishing temperature, which causes nonuniformity of TS in the steel sheet. In the invention, as described above, the amount of Si is decreased to lower the Ar 3 transformation point, so that the finishing temperature can be reliably ensured.
熱間圧延後の一次冷却条件:圧延後の冷却開始時間:2秒以内、平均冷却速度:25℃/秒以上、冷却停止温度:600〜720℃
熱間圧延後、一次冷却開始までに2秒を超える時間放置すると粗大なフェライト粒が生成したり、粗大なTiの炭化物が形成させるため、高強度化や伸びフランジ性の向上を妨げるとともに、鋼板内におけるTSの不均一を増長させる。そのため、圧延後2秒以内に一次冷却を開始する必要がある。また、同様の理由で一次冷却の平均冷却速度は25℃/秒以上とする。
Primary cooling conditions after hot rolling: cooling start time after rolling: within 2 seconds, average cooling rate: 25 ° C / second or more, cooling stop temperature: 600-720 ° C
After hot rolling, if it is left for more than 2 seconds before the start of primary cooling, coarse ferrite grains are formed or coarse Ti carbides are formed, which prevents high strength and stretch flangeability from being improved. Increase the non-uniformity of TS inside. Therefore, it is necessary to start primary cooling within 2 seconds after rolling. For the same reason, the average cooling rate of primary cooling is 25 ° C./second or more.
一次冷却は600〜720℃の温度域で停止させて、引き続く空冷時にフェライト変態と微細なTiの炭化物の析出を促進させる必要がある。冷却停止温度が600℃未満ではフェライト相が十分に生成せず、65%以上の面積率が確保できなくなるとともに、微細なTiの炭化物の密度が低くなる。一方、冷却停止温度が720℃を超えるとフェライト粒やTiの炭化物の粗大化を招き、フェライト相の高強度化が困難になる。したがって、一次冷却の冷却停止温度は600〜720℃とする。 It is necessary to stop the primary cooling in the temperature range of 600 to 720 ° C. and promote the ferrite transformation and the precipitation of fine Ti carbide during the subsequent air cooling. When the cooling stop temperature is less than 600 ° C., the ferrite phase is not sufficiently formed, and an area ratio of 65% or more cannot be secured, and the density of fine Ti carbides is lowered. On the other hand, if the cooling stop temperature exceeds 720 ° C., ferrite grains and Ti carbides become coarse, and it becomes difficult to increase the strength of the ferrite phase. Therefore, the cooling stop temperature of the primary cooling is set to 600 to 720 ° C.
一次冷却後の空冷時間:5〜60秒間
空冷時間が5秒未満ではフェライト相が十分に生成せず、65%以上の面積率が確保できなくなるとともに、微細なTiの炭化物の密度が低くなる。一方、空冷時間が60秒を超えるとフェライト粒やTiの炭化物の粗大化を招き、フェライト相の高強度化が困難になる。したがって、空冷時間は5〜60秒間とする。
Air cooling time after primary cooling: 5 to 60 seconds If the air cooling time is less than 5 seconds, a ferrite phase is not sufficiently formed, and an area ratio of 65% or more cannot be secured, and the density of fine Ti carbides becomes low. On the other hand, if the air cooling time exceeds 60 seconds, ferrite grains and Ti carbides become coarse, and it becomes difficult to increase the strength of the ferrite phase. Therefore, the air cooling time is 5 to 60 seconds.
二次冷却条件:平均冷却速度:25℃/秒以上
熱間圧延後に一次冷却と空冷の組み合わせで得られるフェライト相の面積率65〜80%とフェライト相中の微細なTiの炭化物の析出状態を維持するために、空冷後巻取りまでは25℃/秒以上の平均冷却速度で二次冷却する必要がある。
Secondary cooling condition: Average cooling rate: 25 ° C / second or more The area ratio of the ferrite phase obtained by the combination of primary cooling and air cooling after hot rolling is 65-80% and the precipitation state of fine Ti carbide in the ferrite phase In order to maintain it, it is necessary to perform secondary cooling at an average cooling rate of 25 ° C./second or more after air cooling until winding.
巻取温度:400℃未満
巻取温度が400℃以上となるとマルテンサイト相主体の第二相を生成することが困難となる。したがって、巻取温度は400℃未満とする。材質安定性や板形状安定性を考慮すると、巻取温度は200〜350℃とすることが好ましい。
Winding temperature: less than 400 ° C. When the winding temperature is 400 ° C. or higher, it becomes difficult to produce a second phase mainly composed of a martensite phase. Therefore, the coiling temperature is less than 400 ° C. In consideration of material stability and plate shape stability, the winding temperature is preferably 200 to 350 ° C.
その他の製造条件には通常の条件を適用できる。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、溶融亜鉛系めっき、電気亜鉛系めっき、化成処理を施すことも可能である。ここで、亜鉛系めっきとは、亜鉛および亜鉛を主体とした(すなわち亜鉛を約90%以上含有する)めっきであり、亜鉛のほかにAl、Crなどの合金元素を含んだめっきや亜鉛系めっき後に合金化処理を行っためっきのことである。 Normal conditions can be applied to other manufacturing 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 may be performed, or hot dip galvanizing, electrogalvanizing, or chemical conversion treatment may be performed. Here, zinc-based plating is plating mainly composed of zinc and zinc (that is, containing about 90% or more of zinc), and plating or zinc-based plating containing alloy elements such as Al and Cr in addition to zinc. It is the plating which performed the alloying process later.
表1に示す化学組成とAr3変態点を有する鋼No.A〜Iを転炉で溶製し、連続鋳造法でスラブとした。なお、Ar3変態点は、変態点測定装置を用い、1200℃に10min加熱した後、1℃/秒の冷却速度で冷却する条件により求めた。これらのスラブを、1250℃に加熱し、表2に示す熱延条件で板厚3.2mmのコイル状の熱延鋼板No.1〜21を作製した。そして、酸洗後、コイルの最内周と最外周の一巻き目をカットし、コイルの幅方向の両端部10mmをトリミング後、コイル長手方向に20等分、幅方向に8等分に分割し、端部も含めた189点の位置から圧延方向に平行にJIS 5号引張試験片を採取し、JIS Z 2241に準拠して、クロスヘッド速度10mm/minで引張試験を行い、平均のTSとEl、およびTSの均一性を調べるためにTSのばらつき、すなわち標準偏差ΔTSを求めた。また、189点の位置から穴拡げ試験用試験片を採取し、鉄連規格JFST 1001に準拠して穴拡げ試験を行い、平均のλを求めた。さらに、上記の方法により、フェライト相の面積率およびそのばらつきΔSFやフェライト相とマルテンサイト相の合計の面積率およびフェライト相と第二相のビッカース硬度差ΔHvを求めた。 Steel Nos. A to I having the chemical composition and Ar 3 transformation point shown in Table 1 were melted in a converter and made into a slab by a continuous casting method. The Ar 3 transformation point was determined by using a transformation point measuring device and heating at 1200 ° C. for 10 minutes and then cooling at a cooling rate of 1 ° C./second. These slabs were heated to 1250 ° C., and coiled hot-rolled steel sheets No. 1 to 21 having a thickness of 3.2 mm were produced under the hot-rolling conditions shown in Table 2. Then, after pickling, the innermost and outermost windings of the coil are cut, and both ends 10mm in the width direction of the coil are trimmed, and then divided into 20 equal parts in the coil longitudinal direction and 8 equal parts in the width direction. Then, JIS No. 5 tensile test specimens were collected from 189 points including the end in parallel with the rolling direction, and subjected to a tensile test at a crosshead speed of 10 mm / min in accordance with JIS Z 2241. , El, and TS, in order to examine the uniformity of TS, the standard deviation ΔTS was obtained. In addition, specimens for hole expansion tests were taken from 189 positions and subjected to a hole expansion test in accordance with the Iron Federation standard JFST 1001, to obtain an average λ. Further, the area ratio of the ferrite phase and its variation ΔSF, the total area ratio of the ferrite phase and the martensite phase, and the Vickers hardness difference ΔHv between the ferrite phase and the second phase were determined by the above method.
結果を表3に示す。本発明例では、690〜980MPaのTSが得られ、Elが27%以上、λが50%以上で伸び特性や伸びフランジ性に優れ、かつΔTSが15MPa以下で鋼板内においてTSの均一性に優れていることがわかる。 The results are shown in Table 3. In the example of the present invention, TS of 690 to 980 MPa is obtained, El is 27% or more, λ is 50% or more and excellent in stretch characteristics and stretch flangeability, and ΔTS is 15 MPa or less and excellent in TS uniformity in the steel sheet. You can see that
表1の鋼No.Aの成分組成を有するスラブを、1250℃に加熱し、仕上温度を800〜930℃の範囲に変えて熱間圧延後、1.5秒で一次冷却で開始し、平均冷却速度110℃/秒で700℃まで冷却後、50秒間空冷し、平均冷却速度50℃/秒で二次冷却後、350℃の巻取温度で巻取って板厚3.2mmのコイル状の熱延鋼板を作製した。そして、実施例1と同様な方法でフェライト相と第二相のビッカース硬度差ΔHvを求めた。 The slab having the composition of steel No. A in Table 1 is heated to 1250 ° C, the finishing temperature is changed to the range of 800-930 ° C, hot rolling is started, and primary cooling is started in 1.5 seconds, and the average cooling rate After cooling to 700 ° C at 110 ° C / second, air-cooled for 50 seconds, secondary cooled at an average cooling rate of 50 ° C / second, wound at a coiling temperature of 350 ° C, and coiled hot-rolled steel sheet with a thickness of 3.2 mm Was made. Then, the Vickers hardness difference ΔHv between the ferrite phase and the second phase was determined in the same manner as in Example 1.
図1に仕上温度とΔHvとの関係を示したが、仕上温度を(Ar3変態点＋100℃)以上、すなわちAr3変態点が795℃の鋼No.Aでは仕上温度を895℃以上である900℃以上にすれば、250以下のΔHvが安定して得られることがわかる。 Fig. 1 shows the relationship between the finishing temperature and ΔHv. The finishing temperature is (Ar 3 transformation point + 100 ° C) or higher, that is, the finishing temperature is 895 ° C or higher for Steel No. A with Ar 3 transformation point of 795 ° C. It can be seen that if the temperature is 900 ° C. or higher, a ΔHv of 250 or lower can be stably obtained.
次に、上記において作成した熱延鋼板の一部について、実施例1と同様に、コイルの189点の位置から試験片を採取して、上記の方法でフェライト相の面積率のばらつきΔSFおよびTSのばらつきΔTSを求めた。 Next, for a part of the hot-rolled steel sheet prepared above, as in Example 1, a test piece was taken from the position of 189 points of the coil, and the variation in the area ratio of the ferrite phase ΔSF and TS by the above method The variation ΔTS of was obtained.
図2にフェライト面積率のばらつきΔSFとTSのばらつきΔTSの関係を示す。図2より、ΔSFとΔTSには相関があり、フェライト相の面積率のばらつきΔSFを2%以下とすることにより、TSのばらつきΔTSを15MPa以下にできることがわかる。
FIG. 2 shows the relationship between the ferrite area ratio variation ΔSF and the TS variation ΔTS. From FIG. 2, it can be seen that ΔSF and ΔTS have a correlation, and the variation ΔTS of TS can be made 15 MPa or less by making the variation ΔSF of the ferrite
また、図3に仕上温度とフェライト相の面積率のばらつきΔSFとの関係を示す。図3より、Ar3変態点が795℃の鋼No.Aでは仕上温度を895℃以上である900℃以上にすれば、フェライト相の面積率のばらつきΔSFを2%以下と小さくでき、鋼板内におけるTSの均一化を図れることがわかる。 FIG. 3 shows the relationship between the finishing temperature and the variation ΔSF in the area ratio of the ferrite phase. According to Fig. 3, in steel No. A with an Ar 3 transformation point of 795 ° C, if the finishing temperature is 900 ° C, which is 895 ° C or higher, the variation in the ferrite phase area ratio ΔSF can be reduced to 2% or less, It can be seen that uniform TS can be achieved.
Claims (2)
0.05≦C*≦0.09・・・(1)
0.03≦Ti*≦0.10・・・(2)
ただし、C*=[C]-0.55×Ti*、Ti*=[Ti]-48×[N]/14、[M]は元素Mの含有量(質量%)を表す。 In mass%, C: 0.060 to 0.150%, Si: 0.1% or less, Mn: 0.8 to 1.8%, P: 0.030% or less, S: 0.005% or less, Al: 0.005 to 0.1%, N: 0.005% or less, Ti : The second phase containing 0.032 to 0.120%, the balance consisting of Fe and unavoidable impurities, having the composition satisfying the following formulas (1) and (2), including the ferrite phase and the martensite phase The area ratio of the ferrite phase in the entire structure is 65 to 80%, the total area ratio of the ferrite phase and the martensite phase in the entire structure is 95% or more, the area of the ferrite phase A high-strength hot-rolled steel sheet having a microstructure in which the rate variation ΔSF is 2% or less and the Vickers hardness difference ΔHv between the ferrite phase and the second phase is 250 or less;
0.05 ≦ C * ≦ 0.09 ... (1)
0.03 ≦ Ti * ≦ 0.10 ... (2)
However, C * = [C] −0.55 × Ti * , Ti * = [Ti] −48 × [N] / 14, [M] represents the content (mass%) of the element M.
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