JP5487984B2 - High-strength cold-rolled steel sheet excellent in bendability and manufacturing method thereof - Google Patents
High-strength cold-rolled steel sheet excellent in bendability and manufacturing method thereof Download PDFInfo
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- 239000010960 cold rolled steel Substances 0.000 title claims description 28
- 238000004519 manufacturing process Methods 0.000 title claims description 12
- 229910000831 Steel Inorganic materials 0.000 claims description 81
- 239000010959 steel Substances 0.000 claims description 81
- 229910000859 α-Fe Inorganic materials 0.000 claims description 52
- 238000001816 cooling Methods 0.000 claims description 49
- 229910001563 bainite Inorganic materials 0.000 claims description 40
- 238000005096 rolling process Methods 0.000 claims description 38
- 229910000734 martensite Inorganic materials 0.000 claims description 35
- 238000000137 annealing Methods 0.000 claims description 31
- 229910001566 austenite Inorganic materials 0.000 claims description 20
- 238000005097 cold rolling Methods 0.000 claims description 16
- 239000000203 mixture Substances 0.000 claims description 9
- 239000012535 impurity Substances 0.000 claims description 4
- 230000009466 transformation Effects 0.000 description 19
- 238000005452 bending Methods 0.000 description 17
- 239000013078 crystal Substances 0.000 description 16
- 238000000034 method Methods 0.000 description 15
- 230000008569 process Effects 0.000 description 9
- 230000000717 retained effect Effects 0.000 description 9
- 229910052804 chromium Inorganic materials 0.000 description 8
- 229910052802 copper Inorganic materials 0.000 description 8
- 230000000694 effects Effects 0.000 description 8
- 229910052750 molybdenum Inorganic materials 0.000 description 8
- 229910052759 nickel Inorganic materials 0.000 description 8
- 230000007423 decrease Effects 0.000 description 7
- 229910052720 vanadium Inorganic materials 0.000 description 7
- 229910001562 pearlite Inorganic materials 0.000 description 6
- 230000002411 adverse Effects 0.000 description 5
- 239000000463 material Substances 0.000 description 5
- 238000012360 testing method Methods 0.000 description 5
- 238000004804 winding Methods 0.000 description 5
- 230000015572 biosynthetic process Effects 0.000 description 4
- 238000010438 heat treatment Methods 0.000 description 4
- 230000006872 improvement Effects 0.000 description 4
- XEEYBQQBJWHFJM-UHFFFAOYSA-N iron Substances [Fe] XEEYBQQBJWHFJM-UHFFFAOYSA-N 0.000 description 4
- 238000000465 moulding Methods 0.000 description 4
- 238000001953 recrystallisation Methods 0.000 description 4
- 229910045601 alloy Inorganic materials 0.000 description 3
- 239000000956 alloy Substances 0.000 description 3
- 238000005336 cracking Methods 0.000 description 3
- 238000005098 hot rolling Methods 0.000 description 3
- 229910052748 manganese Inorganic materials 0.000 description 3
- 150000004767 nitrides Chemical class 0.000 description 3
- 230000009467 reduction Effects 0.000 description 3
- 238000009628 steelmaking Methods 0.000 description 3
- VEXZGXHMUGYJMC-UHFFFAOYSA-N Hydrochloric acid Chemical compound Cl VEXZGXHMUGYJMC-UHFFFAOYSA-N 0.000 description 2
- 229910001035 Soft ferrite Inorganic materials 0.000 description 2
- 239000000654 additive Substances 0.000 description 2
- 230000000996 additive effect Effects 0.000 description 2
- 238000007796 conventional method Methods 0.000 description 2
- 230000003111 delayed effect Effects 0.000 description 2
- 238000011156 evaluation Methods 0.000 description 2
- 238000005554 pickling Methods 0.000 description 2
- 238000012545 processing Methods 0.000 description 2
- 229920006395 saturated elastomer Polymers 0.000 description 2
- 238000010583 slow cooling Methods 0.000 description 2
- 239000000126 substance Substances 0.000 description 2
- 238000002441 X-ray diffraction Methods 0.000 description 1
- 230000009471 action Effects 0.000 description 1
- 238000005266 casting Methods 0.000 description 1
- 229910001567 cementite Inorganic materials 0.000 description 1
- 238000006243 chemical reaction Methods 0.000 description 1
- 230000000052 comparative effect Effects 0.000 description 1
- 239000012141 concentrate Substances 0.000 description 1
- 238000009749 continuous casting Methods 0.000 description 1
- 238000012937 correction Methods 0.000 description 1
- 238000005520 cutting process Methods 0.000 description 1
- 238000006477 desulfuration reaction Methods 0.000 description 1
- 230000023556 desulfurization Effects 0.000 description 1
- 238000009792 diffusion process Methods 0.000 description 1
- 239000000446 fuel Substances 0.000 description 1
- 238000000265 homogenisation Methods 0.000 description 1
- 238000010191 image analysis Methods 0.000 description 1
- 230000001771 impaired effect Effects 0.000 description 1
- 229910052742 iron Inorganic materials 0.000 description 1
- KSOKAHYVTMZFBJ-UHFFFAOYSA-N iron;methane Chemical compound C.[Fe].[Fe].[Fe] KSOKAHYVTMZFBJ-UHFFFAOYSA-N 0.000 description 1
- 239000011159 matrix material Substances 0.000 description 1
- 238000005259 measurement Methods 0.000 description 1
- 238000002844 melting Methods 0.000 description 1
- 230000008018 melting Effects 0.000 description 1
- 239000002184 metal Substances 0.000 description 1
- 229910052751 metal Inorganic materials 0.000 description 1
- 150000001247 metal acetylides Chemical class 0.000 description 1
- 239000003595 mist Substances 0.000 description 1
- 230000003287 optical effect Effects 0.000 description 1
- 239000002245 particle Substances 0.000 description 1
- 239000002244 precipitate Substances 0.000 description 1
- 238000003303 reheating Methods 0.000 description 1
- 230000000630 rising effect Effects 0.000 description 1
- 239000006104 solid solution Substances 0.000 description 1
- 230000000087 stabilizing effect Effects 0.000 description 1
- 230000003746 surface roughness Effects 0.000 description 1
- 238000009864 tensile test Methods 0.000 description 1
- XLYOFNOQVPJJNP-UHFFFAOYSA-N water Substances O XLYOFNOQVPJJNP-UHFFFAOYSA-N 0.000 description 1
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- Heat Treatment Of Sheet Steel (AREA)
Description
本発明は、厳しい形状にプレス成形される自動車部品などに供して好適な高強度冷延鋼板およびその製造方法に関し、鋼板の組織がベイナイト相を主体とし、また、加工性や溶接性を阻害するＣの含有量を少なくして、Cu、Ni、Cr、Mo、Ｖなど高価な元素を積極的に含有させることなしに、高い伸び（El）、高い穴拡げ率（λ）のみならず、優れた曲げ性を有し、しかも引張強度（TS）：1180MPa以上という高強度を併せて実現しようとするものである。 The present invention relates to a high-strength cold-rolled steel sheet suitable for use in automobile parts and the like that are press-formed into a strict shape and a method for producing the same. Not only high elongation (El) and high hole expansion ratio (λ) but also excellent without reducing the C content and actively including expensive elements such as Cu, Ni, Cr, Mo, V In addition, it is intended to achieve a high bending strength with a tensile strength (TS) of 1180 MPa or more.
近年、車体軽量化による燃費向上および衝突安全性確保の観点から、自動車部品において高強度鋼板の適用が拡大しつつある。しかしながら、鋼板はその高強度化に伴い加工性が低下する傾向にあり、高強度鋼板を、複雑なプレス成形等に適用した場合、鋼板の割れ発生が一番の問題となる。従って、1180MPa級以上の鋼板をプレス成形する際には、曲げ主体の軽加工とならざるを得ない。また、1180MPa級以上に鋼板を高強度化した場合、ＣやMnなどの添加元素量が増加し、溶接性が著しく低下する場合がある。更に、強度を維持するために、ＣやMnの代替元素として、Cu、Ni、Cr、Mo、Ｖなどの高価な希少元素を積極的に添加する場合がある。従って、Cu、Ni、Cr、Mo、Ｖなどの高価な希少元素を添加せずとも、所望の強度を確保しつつ、低Ｃ量で優れた曲げ性を有する鋼板の開発が望まれていた。 In recent years, the application of high-strength steel sheets in automobile parts has been expanding from the viewpoint of improving fuel efficiency and ensuring collision safety by reducing the weight of the vehicle body. However, the workability of steel sheets tends to decrease with increasing strength, and when the high-strength steel sheets are applied to complicated press forming, cracking of the steel sheets becomes the biggest problem. Therefore, when press forming a steel plate of 1180 MPa class or higher, it is unavoidable to be a light bending process. Further, when the strength of the steel plate is increased to 1180 MPa or higher, the amount of additive elements such as C and Mn increases, and the weldability may be remarkably deteriorated. Furthermore, in order to maintain the strength, an expensive rare element such as Cu, Ni, Cr, Mo, or V may be positively added as an alternative element for C or Mn. Accordingly, it has been desired to develop a steel sheet having a low C content and excellent bendability while ensuring a desired strength without adding expensive rare elements such as Cu, Ni, Cr, Mo, and V.
従来技術として、例えば、特許文献１〜４に、鋼成分や組織の限定、熱延条件、焼鈍条件の最適化により、ベイナイト相を利用する高強度冷延鋼板を得る技術が開示されている。
また、例えば、特許文献５には、鋼板が割れることなく成形できる限界曲げ半径（Ｒ）と板厚（ｔ）の比である限界曲げ半径（Ｒ）／板厚（ｔ）（以下、限界曲げ指数という）が、０〜2.0の値である板厚：1.0mmの鋼板の記載がある。
As conventional techniques, for example, Patent Documents 1 to 4 disclose techniques for obtaining high-strength cold-rolled steel sheets using a bainite phase by limiting steel components and structures, optimizing hot rolling conditions, and annealing conditions.
Further, for example, in Patent Document 5, the limit bending radius (R) / plate thickness (t) which is the ratio of the limit bending radius (R) and the plate thickness (t) that can be formed without cracking the steel plate (hereinafter referred to as limit bending). There is a description of a steel sheet having a thickness of 1.0 mm, which is a value of 0 to 2.0).
しかしながら、特許文献１に記載の鋼板組織では、ベイナイト相の比率が限定されていないため、鋼板の引張強度が十分とは言えなかった。
特許文献２に記載の鋼板は、高価な元素であるMoを必須とする不利があった。
特許文献３に記載の鋼板は、高強度化を図る上でCu、Ni、Cr、Mo、Ｖ等の高価な合金元素を必須とする不利があった。
特許文献４に記載の鋼板は、ベイナイト相が主体であるが、延性に寄与するフェライト相の体積分率が少ないため、延性が不十分であった。
特許文献５に記載の鋼板は、高価な元素であるＶを必須とする不利があり、また、鋼板組成が焼戻しマルテンサイト相を主体とし、かつ延性に寄与するフェライト相の体積分率が少ないため、延性が不十分であった。
However, in the steel sheet structure described in Patent Document 1, since the ratio of the bainite phase is not limited, it cannot be said that the tensile strength of the steel sheet is sufficient.
The steel sheet described in Patent Document 2 has a disadvantage of requiring Mo, which is an expensive element.
The steel sheet described in Patent Document 3 has a disadvantage of requiring expensive alloy elements such as Cu, Ni, Cr, Mo, and V in order to increase the strength.
Although the steel sheet described in Patent Document 4 is mainly composed of a bainite phase, the ductility is insufficient because the volume fraction of the ferrite phase contributing to ductility is small.
The steel sheet described in Patent Document 5 has the disadvantage that V, which is an expensive element, is essential, and the steel sheet composition is mainly composed of a tempered martensite phase, and the volume fraction of the ferrite phase that contributes to ductility is small. The ductility was insufficient.
本発明は、従来技術に見られた上記の問題を有利に解決するものであり、溶接性を阻害するＣおよびAlの含有量を少なくし、しかもCu、Ni、Cr、Mo、Ｖ等の高価な合金元素を含有させない成分系であっても、鋼板組織中のオーステナイトから低温変態生成するベイナイト相、マルテンサイト相および残留オーステナイト相の体積分率を制御し、優れた曲げ性を有する引張強度（TS）が1180MPa以上の高強度冷延鋼板を、その有利な製造方法と共に提供することを目的とする。
なお、本発明において、曲げ性に優れるとは、前記した限界曲げ指数が2.5以下の範囲であることを意味する。
The present invention advantageously solves the above-mentioned problems found in the prior art, reduces the contents of C and Al that impair weldability, and is expensive such as Cu, Ni, Cr, Mo, V, etc. Even if it is a component system that does not contain any alloy elements, it controls the volume fraction of the bainite phase, martensite phase and residual austenite phase that form low temperature transformation from austenite in the steel sheet structure, and has excellent bendability ( The object is to provide a high-strength cold-rolled steel sheet having a TS of 1180 MPa or more together with its advantageous production method.
In the present invention, “excellent bendability” means that the above-mentioned limit bending index is in a range of 2.5 or less.
発明者らは、上記の課題を解決すべく、鋭意検討した。その結果、加工性の観点からＣ含有量を低減し、かつCu、Ni、Cr、Mo、Ｖ等の高価な合金元素の含有量を添加しない成分組成を有する鋼板であっても、ベイナイト相を主体としかつ組織全体の最適化を図ることにより、曲げ性に優れた引張強度（TS）：1180MPa以上の高強度冷延鋼板が得られることを見出した。 The inventors diligently studied to solve the above problems. As a result, even if the steel sheet has a component composition that reduces the C content from the viewpoint of workability and does not add the content of expensive alloy elements such as Cu, Ni, Cr, Mo, and V, the bainite phase is reduced. It was found that a high-strength cold-rolled steel sheet with a tensile strength (TS) of 1180 MPa or more excellent in bendability can be obtained by using the main structure and optimizing the entire structure.
本発明は、上記知見に基づきなされたもので、その要旨構成は以下のとおりである。
（１）質量％で、
Ｃ：0.10％以上0.15％以下、
Si：1.0％以上2.0％以下、
Mn：2.0％以上3.0％以下、
Ｐ：0.030％以下、
Ｓ：0.0050％以下、
Al：0.005％以上0.1％以下、
Ｎ：0.01％以下、
Ti：0.005％以上0.050％以下 および
Ｂ：0.0001％以上0.0050％以下
を含有し、残部がFe及び不可避的不純物からなる成分組成を有し、体積分率で、
ベイナイト相が50％以上70％以下、
フェライト相が20％以上40％以下、
マルテンサイト相が１％以上10％以下および
残留オーステナイト相が５％以下
を含む組織を有し、引張強度：1180MPa以上であることを特徴とする、曲げ性に優れた高強度冷延鋼板。
This invention was made | formed based on the said knowledge, The summary structure is as follows.
(1) In mass%,
C: 0.10% to 0.15%,
Si: 1.0% to 2.0%,
Mn: 2.0% to 3.0%,
P: 0.030% or less,
S: 0.0050% or less,
Al: 0.005% to 0.1%,
N: 0.01% or less,
Ti: 0.005% or more and 0.050% or less and B: 0.0001% or more and 0.0050% or less, with the balance being composed of Fe and inevitable impurities,
The bainite phase is 50% to 70%,
Ferrite phase is 20% to 40%,
A high-strength cold-rolled steel sheet having excellent bendability, having a structure including a martensite phase of 1% to 10% and a residual austenite phase of 5% or less and a tensile strength of 1180 MPa or more.
（２）上記（１）に記載の曲げ性に優れた高強度冷延鋼板の製造方法であって、上記（１）に記載の成分組成からなる鋼スラブを、圧延温度：900℃以上1000℃以下で熱間仕上げ圧延し、該熱間仕上げ圧延後、30℃/秒以上100℃/秒以下の速度で冷却し、350℃以上550℃以下の温度域で巻き取り、ついで酸洗後、圧延率：20％以上50％以下の冷間圧延を施して冷延鋼板としたのち、該冷延鋼板に820℃以上920℃以下の温度域で焼鈍処理を施し、該焼鈍後、冷却速度：10℃/秒以上80℃/秒以下で450℃以上550℃以下の冷却停止温度域まで冷却し、該冷却停止温度域からの温度降下量が０〜100℃の温度範囲に100秒以上1000秒以下滞留させることを特徴とする、曲げ性に優れた高強度冷延鋼板の製造方法。 (2) A method for producing a high-strength cold-rolled steel sheet having excellent bendability as described in (1) above, wherein a steel slab comprising the component composition as described in (1) above is rolled at a temperature of 900 ° C. or higher and 1000 ° C. Hot finish rolling at the following, after the hot finish rolling, cooled at a rate of 30 ° C / second to 100 ° C / second, wound up in a temperature range of 350 ° C to 550 ° C, then pickled, rolled Rate: 20% to 50% cold rolled into a cold-rolled steel sheet, the cold-rolled steel sheet was subjected to an annealing treatment in a temperature range of 820 ° C to 920 ° C, and after the annealing, a cooling rate: 10 Cooling to a cooling stop temperature range of 450 ° C to 550 ° C at a rate of ℃ / second to 80 ° C / second, and a temperature drop from the cooling stop temperature range to a temperature range of 0 to 100 ° C for 100 seconds to 1000 seconds A method for producing a high-strength cold-rolled steel sheet excellent in bendability, characterized by being retained.
本発明によれば、限界曲げ指数が2.5以下という曲げ性に優れ、しかも引張強度が1180MPa以上の高強度冷延鋼板を製造することができる。そのため、本発明により得られる高強度冷延鋼板は、特に厳しい形状にプレス成形される自動車部品として好適である。 According to the present invention, it is possible to produce a high-strength cold-rolled steel sheet having excellent bendability with a limit bending index of 2.5 or less and a tensile strength of 1180 MPa or more. Therefore, the high-strength cold-rolled steel sheet obtained by the present invention is suitable as an automobile part that is press-formed into a particularly severe shape.
発明者らは、高強度冷延鋼板の曲げ性向上に関し、鋭意検討した。その結果、低Ｃ成分かつCu、Ni、Cr、Mo、Ｖ等を含有しない成分とし、前述のとおり、体積分率で、ベイナイト相が50％以上70％以下、フェライト相が20％以上40％以下、マルテンサイト相が１％以上10％以下および残留オーステナイト相が５％以下を含む組織とすることで、曲げ性、伸びおよび伸びフランジ性の向上が顕著となることを見出した。
以下、本発明の成分組成および組織の限定理由について具体的に説明する。なお、鋼板中の元素の含有量の単位は何れも「質量％」であるが、以下、特に断らない限り、単に「％」で示す。
The inventors diligently studied on improving the bendability of a high-strength cold-rolled steel sheet. As a result, it is a low C component and does not contain Cu, Ni, Cr, Mo, V, etc., and as described above, the volume fraction is 50% to 70% for the bainite phase and 20% to 40% for the ferrite phase. In the following, it has been found that the improvement in bendability, elongation, and stretch flangeability becomes remarkable when the structure contains 1% to 10% of martensite phase and 5% or less of retained austenite phase.
Hereinafter, the reasons for limiting the component composition and structure of the present invention will be specifically described. The unit of the element content in the steel sheet is “mass%”, but hereinafter, it is simply indicated by “%” unless otherwise specified.
まず、本発明における鋼の化学成分（組成）の適正範囲およびその限定理由は以下の通りである。
Ｃ：0.10％以上0.15％以下
Ｃ量が0.10％に満たないと、鋼板の強度を確保することが困難となる。一方、Ｃ量が0.15％を超えると、マルテンサイト相、残留オーステナイト相の体積分率が多くなり、過度に高強度化し、成形性、スポット溶接性が低下する。そのため、Ｃ量は0.10％以上0.15％以下の範囲とする。好ましくは、0.11％以上0.14％以下の範囲である。
First, the appropriate range of the chemical component (composition) of steel in the present invention and the reasons for its limitation are as follows.
C: 0.10% or more and 0.15% or less If the amount of C is less than 0.10%, it is difficult to ensure the strength of the steel sheet. On the other hand, when the amount of C exceeds 0.15%, the volume fraction of the martensite phase and the retained austenite phase increases, the strength increases excessively, and the formability and spot weldability deteriorate. Therefore, the C content is in the range of 0.10% to 0.15%. Preferably, it is 0.11% or more and 0.14% or less of range.
Si：1.0％以上2.0％以下
Siは、フェライト安定化元素であり、延性に寄与するフェライト相を生成する上で必須の元素である。また、Siはフェライト相を固溶強化することにより、低温変態相とフェライト相との硬度差を低減するため、伸びフランジ性の向上に寄与する。これらの効果を得るためには、Si量は1.0％以上とする必要がある。一方、Si量が2.0％を超えると、鋼板が脆くなるため、成形時の割れを招来し、化成処理性にも悪影響を及ぼす。そのため、Si量は1.0％以上2.0％以下の範囲とする。好ましくは、1.1％以上1.6％以下の範囲である。
Si: 1.0% to 2.0%
Si is a ferrite stabilizing element, and is an essential element for generating a ferrite phase that contributes to ductility. In addition, Si strengthens the ferrite phase in a solid solution, thereby reducing the hardness difference between the low-temperature transformation phase and the ferrite phase, thereby contributing to the improvement of stretch flangeability. In order to obtain these effects, the Si amount needs to be 1.0% or more. On the other hand, if the Si content exceeds 2.0%, the steel sheet becomes brittle, causing cracks during forming and adversely affecting chemical conversion properties. Therefore, the Si content is in the range of 1.0% to 2.0%. Preferably, it is 1.1% or more and 1.6% or less of range.
Mn：2.0％以上3.0％以下
Mnは、鋼板の高強度化に寄与する元素であり、かかる作用はMn量を2.0％以上とすることで認められる。一方、Mn量が3.0％を超えると、過度に焼入れ性が高まり、所望のフェライト相の確保が困難となり、伸びが低下する。従って、Mn量は2.0％以上3.0％以下の範囲とする。好ましくは2.4％以上2.9％以下の範囲である。
Mn: 2.0% to 3.0%
Mn is an element that contributes to increasing the strength of the steel sheet, and this effect is recognized when the Mn content is 2.0% or more. On the other hand, if the amount of Mn exceeds 3.0%, the hardenability is excessively increased, it becomes difficult to secure a desired ferrite phase, and the elongation decreases. Therefore, the Mn content is in the range of 2.0% to 3.0%. Preferably it is 2.4 to 2.9% of range.
Ｐ：0.030％以下
Ｐは、スポット溶接性に悪影響をおよぼすため、本発明においては極力低減することが好ましいが、0.030％までは許容できる。なお、Ｐ量を過度に低減することは製鋼工程での生産効率の低下を招来し、高コストとなるため、Ｐ量の下限は0.001％程度とすることが好ましい。
P: 0.030% or less P has an adverse effect on spot weldability. Therefore, it is preferably reduced as much as possible in the present invention, but up to 0.030% is acceptable. In addition, excessively reducing the amount of P causes a reduction in production efficiency in the steel making process, resulting in high costs. Therefore, the lower limit of the amount of P is preferably about 0.001%.
Ｓ：0.0050％以下
Ｓは、鋼中でMnSを形成して板状介在物となり、伸びフランジ性を低下させる。そのため、本発明においてＳ量は極力低減することが好ましいが、Ｓ量が0.0050％以下であれば上記問題が顕在化することはない。好ましくは0.0030％以下である。
なお、Ｓ量の過度の低減は、製鋼工程における脱硫コストの増加を招くので、Ｓ量の下限は0.0001％程度とすることが好ましい。
S: 0.0050% or less S forms plate-like inclusions by forming MnS in steel and reduces stretch flangeability. Therefore, in the present invention, it is preferable to reduce the S amount as much as possible, but the above problem does not become apparent if the S amount is 0.0050% or less. Preferably it is 0.0030% or less.
In addition, excessive reduction of the amount of S causes an increase in the desulfurization cost in the steel making process, so the lower limit of the amount of S is preferably about 0.0001%.
Al：0.005％以上0.1％以下
Alは、脱酸剤として使用される。脱酸作用を得るためにはAl量を0.005％以上とすることが必要となるが、Al量が0.1％を超えると、溶接性に悪影響を及ぼす。従って、Al量は0.005％以上0.1％以下の範囲とする。好ましくは0.02％以上0.06％以下の範囲である。
Al: 0.005% to 0.1%
Al is used as a deoxidizer. In order to obtain a deoxidizing action, the Al amount needs to be 0.005% or more. However, if the Al amount exceeds 0.1%, weldability is adversely affected. Therefore, the Al content is in the range of 0.005% to 0.1%. Preferably it is 0.02% or more and 0.06% or less of range.
Ｎ：0.01％以下
本発明において、Ｎは不純物であり極力低減することが好ましいが、0.01％までは許容できる。好ましくは0.0050％以下である。
なお、Ｎ量を過度に低減すると、製鋼工程における脱窒コストの増加を招くので、Ｎ量の下限は0.0001％程度とすることが好ましい。より好ましくは0.0010％以上である。
N: 0.01% or less In the present invention, N is an impurity and is preferably reduced as much as possible, but up to 0.01% is acceptable. Preferably it is 0.0050% or less.
Note that, if the amount of N is excessively reduced, denitrification costs in the steelmaking process are increased, so the lower limit of the amount of N is preferably about 0.0001%. More preferably, it is 0.0010% or more.
Ti：0.005％以上0.050％以下
Tiは、鋼中でＣやＮと結合して微細炭化物や微細窒化物を形成することにより、加熱時の結晶粒の粗大化を抑制し、熱延板組織ならびに焼鈍後の鋼板組織の細粒均一化に有効に作用する。また、窒化物形成によりＢ窒化物の形成を抑制し、後述するＢ添加による焼入れ性を確保する上でも有効である。これらの効果を発現すべくTi量は0.005％以上とすることを要するが、Ti量が0.050％を超えるとこれらの効果は飽和する傾向にある。また、Tiを過度に含有すると、フェライト相中にTiの析出物が過剰に生成し、フェライト相の延性低下、さらには、鋼板の硬質化により熱間圧延時および冷間圧延時の圧延荷重が増大する。従って、Ti量は0.005％以上0.050％以下の範囲とする。好ましくは0.010％以上0.040％以下の範囲である。
Ti: 0.005% to 0.050%
Ti combines with C and N in steel to form fine carbides and fine nitrides, thereby suppressing coarsening of crystal grains during heating, and fine graining of hot rolled sheet structure and steel sheet structure after annealing. It works effectively on homogenization. Further, it is effective in suppressing the formation of B nitride by forming the nitride and ensuring hardenability by adding B, which will be described later. In order to exhibit these effects, the Ti content needs to be 0.005% or more, but when the Ti content exceeds 0.050%, these effects tend to be saturated. In addition, if Ti is contained excessively, Ti precipitates are excessively generated in the ferrite phase, the ductility of the ferrite phase is lowered, and further, the rolling load at the time of hot rolling and cold rolling is increased due to the hardening of the steel sheet. Increase. Accordingly, the Ti content is in the range of 0.005% to 0.050%. Preferably it is 0.010% or more and 0.040% or less of range.
Ｂ:0.0001％以上0.0050％以下
Ｂは、焼入れ性を高め、焼鈍後の冷却過程におけるフェライト相の過剰生成を抑制し、所望のベイナイト相量、およびマルテンサイト相量を得るのに寄与する。かかる効果を得るためには、Ｂ量を0.0001％以上とする必要がある。一方、Ｂ量が0.0050％を超えると上記効果は飽和する。従って、Ｂ量は0.0001％以上0.0050％以下の範囲とする。好ましくは0.0005％以上0.0020％以下の範囲である。
なお、本発明の鋼板において、上記以外の成分はFeおよび不可避的不純物である。ただし、本発明の効果を損なわない範囲内であれば、上記以外の成分の含有を拒むものではない。
B: 0.0001% or more and 0.0050% or less B enhances hardenability, suppresses excessive formation of ferrite phase in the cooling process after annealing, and contributes to obtaining a desired bainite phase amount and martensite phase amount. In order to obtain such an effect, the B amount needs to be 0.0001% or more. On the other hand, when the amount of B exceeds 0.0050%, the above effect is saturated. Accordingly, the B content is in the range of 0.0001% to 0.0050%. Preferably it is 0.0005% or more and 0.0020% or less of range.
In the steel sheet of the present invention, components other than those described above are Fe and inevitable impurities. However, as long as the effects of the present invention are not impaired, the inclusion of components other than those described above is not rejected.
次に、本発明にとって重要な要件の一つである鋼の組織の限定範囲およびその限定理由について説明する。
ベイナイト相の体積分率：50％以上70％以下
オーステナイト相からの低温変態相であるベイナイト相はマルテンサイト相よりも軟質であり、フェライト相との硬度差は、マルテンサイト相よりも小さく、鋼板の成形時には鋼板全体が均一に伸びる作用を助けるため伸びフランジ性の向上に有利である。また、曲げ成形時、鋼板の最表層においても均一に材料が変形するため、曲げ性に優れる。マルテンサイト相より軟質なベイナイト相主体で引張強度を確保するためには、50％以上の体積分率のベイナイト相が必要となる。50％に満たない場合、所望の引張強度が確保できないかまたは硬質なマルテンサイト相が増えることとなり、伸びフランジ性の低下につながる。一方、ベイナイト相が体積分率で70％を超えると、延性に寄与する所定量のフェライト相の確保が困難となり、鋼板が過度に硬質化し、伸びの確保が困難となる。従って、ベイナイト相は体積分率で50％以上70％以下の範囲とする。好ましくは50％以上68％以下の範囲である。
Next, the limited range of the steel structure, which is one of the important requirements for the present invention, and the reason for the limitation will be described.
The volume fraction of the bainite phase: 50% or more and 70% or less The bainite phase, which is a low-temperature transformation phase from the austenite phase, is softer than the martensite phase, and the hardness difference from the ferrite phase is smaller than the martensite phase. This is advantageous for improving the stretch flangeability because it helps the entire steel sheet to be stretched uniformly during the forming of. Moreover, since the material is uniformly deformed even in the outermost layer of the steel plate during bending, the bendability is excellent. A bainite phase having a volume fraction of 50% or more is required to ensure tensile strength mainly composed of a bainite phase that is softer than the martensite phase. If it is less than 50%, the desired tensile strength cannot be ensured or the hard martensite phase increases, leading to a decrease in stretch flangeability. On the other hand, when the bainite phase exceeds 70% in volume fraction, it becomes difficult to secure a predetermined amount of ferrite phase that contributes to ductility, the steel sheet becomes excessively hard, and it becomes difficult to ensure elongation. Therefore, the bainite phase is in the range of 50% to 70% in terms of volume fraction. Preferably, it is in the range of 50% or more and 68% or less.
ベイナイト相の平均結晶粒径については、特に限定されるものではないが、ベイナイト相の平均結晶粒径が1μmに満たない場合、フェライト相より硬質なベイナイト相が微細に分散して存在するため、加工時の変形能におよぼすベイナイト相の寄与が大きくなり、また、細粒化に起因して過度に高強度化し、優れた成形性を確保することが困難となる。一方、ベイナイト相の平均結晶粒径が10μmを超えて過度に粗大化すると不均一な組織となり、成形時の均一な変形を阻害する要因となり優れた成形性を確保することが困難となる。従って、ベイナイト相の平均結晶粒径は1μm以上10μm以下の範囲が好ましい。 The average crystal grain size of the bainite phase is not particularly limited, but when the average crystal grain size of the bainite phase is less than 1 μm, the harder bainite phase than the ferrite phase is finely dispersed, The contribution of the bainite phase to the deformability at the time of processing becomes large, and it becomes difficult to ensure excellent moldability by increasing the strength excessively due to the finer grain. On the other hand, when the average crystal grain size of the bainite phase exceeds 10 μm and becomes excessively coarse, a non-uniform structure is formed, and it becomes a factor that inhibits uniform deformation during molding, so that it is difficult to ensure excellent moldability. Accordingly, the average crystal grain size of the bainite phase is preferably in the range of 1 μm to 10 μm.
フェライト相の体積分率：20％以上40％以下
軟質なフェライト相は鋼板の伸び向上に寄与するため、本発明では、フェライト相を体積分率で20％以上とする必要がある。一方、フェライト相が体積分率で40％を超えると、低温変態相との組み合わせにもよるが引張強度：1180MPaの確保が困難となる場合がある。従って、フェライト相は体積分率で20％以上40％以下の範囲とする。
Volume fraction of ferrite phase: 20% or more and 40% or less Since the soft ferrite phase contributes to the improvement of elongation of the steel sheet, the ferrite phase needs to be 20% or more in terms of volume fraction in the present invention. On the other hand, if the ferrite phase exceeds 40% in volume fraction, it may be difficult to ensure a tensile strength of 1180 MPa, depending on the combination with the low temperature transformation phase. Therefore, the ferrite phase is in the range of 20% to 40% in terms of volume fraction.
フェライト相の平均結晶粒径については、特に限定されるものではないが、フェライト相の平均結晶粒径が1μmに満たない場合、硬質な低温変態相（ベイナイト相およびマルテンサイト相）が近接して存在し、加工時の変形能におよぼす低温変態相の寄与が大きくなるため、優れた成形性を確保することが困難となる。一方、フェライト相の平均結晶粒径が10μmを超えて過度に粗大化すると不均一な組織となるため、成形時に不均一な変形が生じ、優れた成形性を確保することが困難となる。従って、フェライト相の平均結晶粒径は1μm以上10μm以下の範囲が好ましい。 The average crystal grain size of the ferrite phase is not particularly limited, but when the average crystal grain size of the ferrite phase is less than 1 μm, hard low-temperature transformation phases (bainite phase and martensite phase) are close to each other. Since the contribution of the low temperature transformation phase to the deformability during processing increases, it becomes difficult to ensure excellent moldability. On the other hand, if the average crystal grain size of the ferrite phase exceeds 10 μm and becomes too coarse, a non-uniform structure is formed. Therefore, non-uniform deformation occurs during molding, and it becomes difficult to ensure excellent moldability. Accordingly, the average crystal grain size of the ferrite phase is preferably in the range of 1 μm to 10 μm.
マルテンサイト相の体積分率：１％以上10％以下
マルテンサイト相を、体積分率で１％以上10％以下の範囲内とする組織とすることで、強度と加工性（曲げ性、伸び、伸びフランジ性）との良好な材質バランスが得られる。マルテンサイト相が体積分率で１％未満の場合、引張強度（TS）：1180MPaの確保が困難となるか、または引張強度を確保するためにベイナイト相の体積分率を増やす必要があるが、この時は、鋼板の成形性が低下する。一方、マルテンサイト相が体積分率で10％超の場合、鋼板が過度に高強度化し、伸びが著しく低下し、またフェライト相との硬度差に起因して優れた伸びフランジ性の確保が困難となる。従って、マルテンサイト相は体積分率で１％以上10％以下の範囲とする。
Volume fraction of martensite phase: 1% or more and 10% or less Strength and workability (bendability, elongation, etc.) by making the martensite phase within the range of 1% or more and 10% or less in volume fraction. Good material balance with stretch flangeability). If the martensite phase is less than 1% in volume fraction, it will be difficult to secure tensile strength (TS): 1180 MPa, or it will be necessary to increase the volume fraction of bainite phase to ensure tensile strength. At this time, the formability of the steel sheet decreases. On the other hand, if the martensite phase is more than 10% in volume fraction, the steel sheet becomes excessively strong, the elongation is remarkably reduced, and it is difficult to ensure excellent stretch flangeability due to the hardness difference from the ferrite phase. It becomes. Therefore, the martensite phase is in the range of 1% to 10% in terms of volume fraction.
マルテンサイト相の平均結晶粒径については、特に限定されるものではないが、マルテンサイト相の平均結晶粒径が0.5μmより小さい場合、硬質なマルテンサイト相がフェライト母相中に微細に分散、またはベイナイト相に隣接するため、成形時にフェライト相またはベイナイト相とマルテンサイト相の界面におけるボイドの発生が多くなり、伸びや伸びフランジ性が低下する。また、細粒化に起因して過度に高強度化し、優れた成形性を確保することが困難となる。
一方、マルテンサイト相の平均結晶粒径が５μmを超えて過度に粗大化すると、不均一な組織となり、成形時の均一な変形を阻害する要因となり優れた成形性を確保することが困難となる。従って、マルテンサイト相の平均結晶粒径は0.5μm以上５μm以下の範囲が好ましい。
The average grain size of the martensite phase is not particularly limited, but if the average grain size of the martensite phase is smaller than 0.5 μm, the hard martensite phase is finely dispersed in the ferrite matrix, Or, since it is adjacent to the bainite phase, voids are often generated at the interface between the ferrite phase or the bainite phase and the martensite phase at the time of molding, and elongation and stretch flangeability are deteriorated. In addition, it becomes difficult to secure an excellent moldability by excessively increasing the strength due to the finer particles.
On the other hand, when the average crystal grain size of the martensite phase is excessively coarsened exceeding 5 μm, a non-uniform structure is formed, and it becomes difficult to ensure excellent moldability because it inhibits uniform deformation during molding. . Therefore, the average crystal grain size of the martensite phase is preferably in the range of 0.5 μm to 5 μm.
残留オーステナイト相の体積分率：５％以下
残留オーステナイト相は、歪誘起変態により延性の向上に寄与するため、体積分率で２％以上存在しているほうが好ましいが、残留オーステナイト相はＣ濃度が高く、硬質相であるため、体積分率で５％を超えて鋼板中に存在すると、伸びフランジ性に悪影響を及ぼす。また、その場合、鋼板上に、局所的に硬質な部分が存在することとなり、曲げ成形時の材料の均一な変形を阻害する要因となるため、優れた曲げ性を確保することが難しくなる。すなわち、曲げ性および伸びフランジ性の確保の観点から残留オーステナイト相は少ないほうが好ましく、本発明では、残留オーステナイト相の体積分率は５％以下とする。
Volume fraction of retained austenite phase: 5% or less The retained austenite phase is preferably present in a volume fraction of 2% or more because it contributes to improvement of ductility by strain-induced transformation. Since it is high and is a hard phase, if it exceeds 5% in the volume fraction, it will adversely affect stretch flangeability. Further, in that case, a locally hard portion exists on the steel plate, which becomes a factor that hinders uniform deformation of the material at the time of bending, so that it is difficult to ensure excellent bendability. That is, from the viewpoint of securing bendability and stretch flangeability, it is preferable that the retained austenite phase is small. In the present invention, the volume fraction of the retained austenite phase is 5% or less.
次に本発明の高強度冷延鋼板の製造方法について説明する。
上記の成分組成を有する鋼スラブを、圧延温度：900℃以上1000℃以下で熱間仕上げ圧延し、該熱間仕上げ圧延後の冷却速度を30℃/秒以上100℃/秒以下、巻き取り温度を350℃以上550℃以下とし、酸洗した後、圧延率：20％以上50％以下の冷間圧延を施して冷延鋼板とし、ついで、該冷延鋼板に820℃以上920℃以下の温度域で焼鈍処理を施し、冷却速度：10℃/秒以上80℃/秒以下で450℃以上550℃以下の冷却停止温度域まで冷却して、該冷却停止温度からの温度降下量が０〜100℃の温度範囲に100秒以上1000秒以下滞留させる。かかる製造方法により本発明の目的とする高強度冷延鋼板が得られるが、鋼板にスキンパス圧延を施しても良い。
以下、製造条件の適正範囲およびその限定理由について説明する。
Next, the manufacturing method of the high-strength cold-rolled steel sheet of this invention is demonstrated.
A steel slab having the above component composition is hot finish rolled at a rolling temperature of 900 ° C. or higher and 1000 ° C. or lower, and a cooling rate after the hot finish rolling is 30 ° C./second or higher and 100 ° C./second or lower, a winding temperature. 350 ° C. or more and 550 ° C. or less, pickling, and cold rolling at a rolling rate of 20% or more and 50% or less to obtain a cold-rolled steel plate, and then the temperature of 820 ° C. or more and 920 ° C. or less to the cold-rolled steel plate The cooling rate is 10 ° C./second or more and 80 ° C./second or less to a cooling stop temperature region of 450 ° C. or more and 550 ° C. or less, and the temperature drop from the cooling stop temperature is 0 to 100 Remain in the temperature range of ℃ 100 to 1000 seconds. Although the high-strength cold-rolled steel sheet that is the object of the present invention is obtained by such a manufacturing method, the steel sheet may be subjected to skin pass rolling.
Hereinafter, the appropriate range of manufacturing conditions and the reason for limitation will be described.
本発明において、熱間仕上げ圧延前の工程に関しては常法に従って行えばよく、例えば、上記の成分組成範囲に調製した鋼を溶製、鋳造して得られた鋼スラブを用いることができる。また、本発明においては、連続鋳造スラブ、造塊−分塊スラブは勿論のこと、厚み：50〜100mm程度の薄スラブを用いることができ、特に薄スラブの場合は、再加熱なしに直接熱間圧延工程に供することができる。 In the present invention, the process before hot finish rolling may be performed in accordance with a conventional method. For example, a steel slab obtained by melting and casting steel prepared in the above component composition range can be used. In the present invention, not only continuous casting slabs and ingot-splitting slabs, but also thin slabs with a thickness of about 50 to 100 mm can be used. Especially in the case of thin slabs, direct heating without reheating is possible. It can use for a hot rolling process.
熱間仕上げ圧延温度：900℃以上1000℃以下
熱間仕上げ圧延温度が900℃より低い場合、熱間仕上げ圧延後の組織がバンド状の展伸粒組織となり、冷延焼鈍後もバンド状の展伸粒組織のままである。そのため、伸びフランジ性が低下する。また、熱延板の結晶粒径が微細化し、冷延焼鈍後の冷却過程において、鋼板にフェライト相が生成しやすくなるため、引張強度：1180 MPa以上の確保が困難となる。一方、1000℃を超えると、熱間仕上げ圧延後の組織が粗大粒となり、冷延焼鈍後の組織も粗大なままである。そのため、冷延焼鈍後の冷却中のフェライト相の生成が遅延することとなり、過度に硬度化すると共に、伸びフランジ性が低下する傾向を示す。また、この場合、熱間仕上げ圧延後に高温で滞留することとなるため、スケール厚が厚くなって、酸洗後の表面の凹凸が大きくなり、冷延焼鈍後の鋼板の曲げ性に悪影響を及ぼす結果となる。従って、熱間仕上げ圧延温度は900℃以上1000℃以下とする。
Hot finish rolling temperature: 900 ° C or more and 1000 ° C or less When the hot finish rolling temperature is lower than 900 ° C, the structure after hot finish rolling becomes a band-like expanded grain structure, and the band-like extension after cold rolling annealing. The stretched structure remains. Therefore, stretch flangeability falls. In addition, since the crystal grain size of the hot-rolled sheet becomes finer and a ferrite phase is easily generated in the steel sheet in the cooling process after cold-rolling annealing, it is difficult to ensure a tensile strength of 1180 MPa or more. On the other hand, when it exceeds 1000 ° C., the structure after hot finish rolling becomes coarse, and the structure after cold rolling annealing also remains coarse. Therefore, the production | generation of the ferrite phase in the cooling after cold rolling annealing will be delayed, and while it hardens excessively, it shows the tendency for stretch flangeability to fall. Moreover, in this case, since it will stay at a high temperature after hot finish rolling, the scale thickness becomes thick, the surface unevenness after pickling becomes large, and the bendability of the steel sheet after cold rolling annealing is adversely affected. Result. Therefore, the hot finish rolling temperature is set to 900 ° C. or higher and 1000 ° C. or lower.
熱間仕上げ圧延後の冷却速度：30℃/秒以上100℃/秒以下
熱間仕上げ圧延後の冷却速度が30℃/秒より遅い場合、フェライト相の体積分率が多くなるとともに、鋼板組織のパーライト変態が促進され、熱間仕上げ圧延後の組織がフェライト相とパーライト相のバンド状組織となり、冷延焼鈍後もそのままバンド状組織が維持され、鋼板の伸びフランジ性が劣化する。一方、熱間仕上げ圧延後の冷却速度が、過度に速すぎると、所望の温度で鋼板の冷却を停止することが困難となり、また、コスト的にも不利となるため、上限は100℃/秒とする。従って、熱間仕上げ圧延後の冷却速度は30℃/秒以上100℃/秒以下とする。
Cooling rate after hot finish rolling: 30 ° C / second or more and 100 ° C / second or less When the cooling rate after hot finish rolling is slower than 30 ° C / second, the volume fraction of the ferrite phase increases and the steel structure The pearlite transformation is promoted, the structure after hot finish rolling becomes a band-like structure of a ferrite phase and a pearlite phase, the band-like structure is maintained as it is after cold rolling annealing, and the stretch flangeability of the steel sheet deteriorates. On the other hand, if the cooling rate after hot finish rolling is too high, it becomes difficult to stop the cooling of the steel sheet at a desired temperature, and it is also disadvantageous in terms of cost, so the upper limit is 100 ° C / sec. And Therefore, the cooling rate after hot finish rolling is 30 ° C./second or more and 100 ° C./second or less.
巻き取り温度：350℃以上550℃以下
巻き取り温度が550℃を超えると、熱間仕上げ圧延後の組織は、フェライト相の体積分率が多くなるとともに、フェライト相とパーライト相の混在した組織となる。この組織は、Ｃ濃度の低いフェライト相の領域とＣ濃度の高いパーライト相の領域とが存在している不均一な組織であり、連続焼鈍のような短時間の熱処理では冷延焼鈍後も不均一な組織のままであり、鋼板の曲げ性、伸びフランジ性が共に劣化する。一方、巻き取り温度が過度に低すぎると、コスト的に不利となり、また、鋼板が過度に硬質化して冷間圧延時の変形抵抗が増大するため、冷間圧延性が低下する。従って、巻き取り温度は350℃以上550℃以下とする。
Winding temperature: 350 ° C or more and 550 ° C or less When the winding temperature exceeds 550 ° C, the structure after hot finish rolling increases the volume fraction of the ferrite phase and the structure in which the ferrite phase and pearlite phase are mixed. Become. This structure is a heterogeneous structure in which a ferrite phase region having a low C concentration and a pearlite phase region having a high C concentration exist, and a short-time heat treatment such as continuous annealing is not effective even after cold rolling annealing. The uniform structure remains, and both the bendability and stretch flangeability of the steel sheet deteriorate. On the other hand, when the coiling temperature is too low, it is disadvantageous in terms of cost, and the steel sheet becomes excessively hard and deformation resistance at the time of cold rolling is increased, so that cold rolling property is lowered. Accordingly, the winding temperature is set to 350 ° C. or higher and 550 ° C. or lower.
冷間圧延率：20％以上50％以下
圧延率が20％に満たないと、鋼板中に歪が均一に導入されないため、鋼板中で再結晶の進み具合にバラツキが生じ、粗大な粒と微細な粒が存在する不均一な組織となり、伸びフランジ性が劣化する。また、冷間圧延後の焼鈍過程における再結晶、変態挙動が遅延し、焼鈍中のオーステナイト相の量が減少するため、最終的に得られる鋼板中のフェライト相の量が過剰となる。その結果、鋼板の引張強度は低下する。一方、圧延率が50％を超えると、再結晶が急速に進み、粒成長が促進されるため、結晶粒径が粗大化する。また、冷却中のフェライト相の生成が抑制され過度に硬度化し、曲げ性、伸びフランジ性が劣化する。よって冷間圧延率は20％以上50％以下の範囲とする。
Cold rolling rate: 20% or more and 50% or less If the rolling rate is less than 20%, strain is not uniformly introduced into the steel sheet, resulting in variations in the progress of recrystallization in the steel sheet, resulting in coarse grains and fineness. As a result, a non-uniform structure in which fine grains are present is formed, and stretch flangeability is deteriorated. Moreover, since the recrystallization and transformation behavior in the annealing process after cold rolling is delayed and the amount of austenite phase during annealing is reduced, the amount of ferrite phase in the finally obtained steel sheet becomes excessive. As a result, the tensile strength of the steel sheet decreases. On the other hand, when the rolling rate exceeds 50%, recrystallization proceeds rapidly and grain growth is promoted, so that the crystal grain size becomes coarse. Further, the formation of a ferrite phase during cooling is suppressed, the hardness is excessively increased, and the bendability and stretch flangeability deteriorate. Therefore, the cold rolling rate is in the range of 20% to 50%.
焼鈍温度：820℃以上920℃以下
焼鈍温度が820℃に満たない場合、加熱焼鈍中のフェライト分率が高まることに起因して、焼鈍後に最終的に得られるフェライト相の体積分率が過剰となり、引張強度：1180MPa以上の確保が困難となる。また、ＣやMnなどの添加元素の拡散が不十分な状態である濃度ムラが発生して、鋼板組織が低温変態相の偏在する不均一な組織となり、鋼板の加工性（曲げ性、伸び、伸びフランジ性）が劣化する傾向を示す。一方、920℃を超えた場合、オーステナイト単相の温度域まで加熱すると、オーステナイト粒径が過度に粗大化し、その後の冷却過程で生成するフェライト相の量が減少し、伸びが低下する。また、フェライト相や低温変態相の結晶粒径が粗大化し、伸びフランジ性が劣化する。従って、焼鈍温度は820℃以上920℃以下の範囲とする。より好ましくは、830℃以上920℃以下の範囲である。
Annealing temperature: 820 ° C or more and 920 ° C or less If the annealing temperature is less than 820 ° C, the volume fraction of the ferrite phase finally obtained after annealing becomes excessive due to the increase of the ferrite fraction during heat annealing. , Tensile strength: It is difficult to ensure 1180 MPa or more. Further, density unevenness in which the diffusion of additive elements such as C and Mn is insufficient occurs, and the steel sheet structure becomes a non-uniform structure in which low-temperature transformation phases are unevenly distributed, and the workability (bendability, elongation, Elongation flangeability) tends to deteriorate. On the other hand, when the temperature exceeds 920 ° C., heating to the temperature range of the austenite single phase excessively coarsens the austenite grain size, reduces the amount of ferrite phase generated in the subsequent cooling process, and decreases elongation. Further, the crystal grain size of the ferrite phase and the low-temperature transformation phase becomes coarse, and the stretch flangeability deteriorates. Accordingly, the annealing temperature is in the range of 820 ° C. or more and 920 ° C. or less. More preferably, it is the range of 830 degreeC or more and 920 degrees C or less.
冷却速度：10℃/秒以上80℃/秒以下
焼鈍処理後の冷却速度は、軟質なフェライト相と硬質な低温変態相（ベイナイト相およびマルテンサイト相）の体積分率を制御し、1180MPa級以上の引張強度と、優れた曲げ性を確保する上で重要である。
平均冷却速度が80℃/秒を超えると、冷却中のフェライト相生成が抑制され、低温変態相であるベイナイト相およびマルテンサイト相の体積分率が増加するため、1180MPa級以上の引張強度の確保は容易となるが、曲げ性および延性は劣化する。一方、10℃/秒未満であると、冷却過程中に生成するフェライト相の量が過剰となり、1180MPa級以上の引張強度の確保が困難となる。従って、焼鈍処理後の冷却速度は10℃/秒以上80℃/秒以下の範囲とする。好ましくは10℃/秒以上45℃/秒以下の範囲である。
なお、上記冷却速度に制御すべく、焼鈍処理後の鋼板の冷却はガス冷却とすることが好ましいが、その他、炉冷、ミスト冷却、ロール冷却および水冷などの方法を用いることができ、またはそれらを組み合わせて使用することも可能である。
Cooling rate: 10 ° C / second or more and 80 ° C / second or less The cooling rate after annealing treatment controls the volume fraction of the soft ferrite phase and the hard low-temperature transformation phase (bainite phase and martensite phase), and is over 1180 MPa class It is important to ensure the tensile strength and excellent bendability of the steel.
If the average cooling rate exceeds 80 ° C / sec, ferrite phase formation during cooling is suppressed and the volume fraction of the low-temperature transformation phase, bainite phase and martensite phase, increases, ensuring a tensile strength of 1180 MPa class or higher. Is easy, but the bendability and ductility deteriorate. On the other hand, if it is less than 10 ° C./second, the amount of ferrite phase generated during the cooling process becomes excessive, and it becomes difficult to secure a tensile strength of 1180 MPa class or higher. Therefore, the cooling rate after the annealing treatment is set in the range of 10 ° C./second to 80 ° C./second. The range is preferably 10 ° C./second or more and 45 ° C./second or less.
In order to control the cooling rate, the steel plate after annealing is preferably cooled by gas, but other methods such as furnace cooling, mist cooling, roll cooling, and water cooling can be used, or It is also possible to use in combination.
冷却停止温度：450℃以上550℃以下
冷却停止温度を450℃未満とすると、低温まで冷却されることとなり、フェライト相の生成量が過剰に増加し、1180MPa級以上の引張強度の確保が困難となる。一方、冷却停止温度が550℃を超えると、低温変態相のベイナイトの体積分率が多くなりすぎ、また、パーライト相も併せて生成するために、むしろ鋼板の曲げ性および伸びフランジ性は低下する。
以上の理由により、本発明においては、ベイナイト相を主体とし、フェライト相の体積分率を制御し、1180MPa級以上の引張強度を確保するとともに優れた延性、伸びフランジ性および曲げ性をバランス良く得るのに、冷却停止温度は450℃以上550℃以下の範囲とする。
Cooling stop temperature: 450 ° C or higher and 550 ° C or lower If the cooling stop temperature is lower than 450 ° C, it will be cooled to a low temperature, the amount of ferrite phase generated will increase excessively, and it will be difficult to secure a tensile strength of 1180 MPa class or higher. Become. On the other hand, when the cooling stop temperature exceeds 550 ° C., the volume fraction of the low-temperature transformation phase bainite increases too much, and the pearlite phase is also generated, so the bendability and stretch flangeability of the steel sheet are rather lowered. .
For the above reasons, in the present invention, the bainite phase is the main component, the volume fraction of the ferrite phase is controlled, the tensile strength of 1180 MPa class or higher is ensured, and excellent ductility, stretch flangeability and bendability are obtained in a well-balanced manner. However, the cooling stop temperature is in the range of 450 ° C to 550 ° C.
鋼板滞留温度：冷却停止温度からの温度降下量が０〜100℃の温度範囲
オーステナイト相からの低温変態相は、変態温度が低いほど硬くなる。そのため、冷却停止後の鋼板滞留温度は、ベイナイト相の強度を制御をするのに重要である。
オーステナイト相からの低温変態相は、等温保持、すなわち温度降下量が０℃の場合、強度が均一なベイナイト相が主体の組織となる。そのため、曲げおよび伸びフランジ成形を行った時に、鋼板の均一な変形が可能となり、優れた成形性を示す。一方、滞留中の鋼板温度が100℃を超えて降下すると、マルテンサイト相が過剰に生成し、不均一な組織となる。そのため、鋼板を成形する際、フェライト相とマルテンサイト相の粒界に歪が局所的に集中し、優れた曲げ性および伸びフランジ性を達成することが困難となる。従って、鋼板の冷却停止温度からの温度降下量は０〜100℃の範囲とする。
Steel plate residence temperature: Temperature range in which the temperature drop from the cooling stop temperature is 0 to 100 ° C. The low temperature transformation phase from the austenite phase becomes harder as the transformation temperature is lower. Therefore, the steel plate residence temperature after cooling is stopped is important for controlling the strength of the bainite phase.
The low temperature transformation phase from the austenite phase is isothermally maintained, that is, when the temperature drop is 0 ° C., the bainite phase having a uniform strength is the main structure. Therefore, when bending and stretch flange forming are performed, uniform deformation of the steel sheet is possible, and excellent formability is exhibited. On the other hand, when the temperature of the staying steel plate falls below 100 ° C., the martensite phase is excessively generated, resulting in a non-uniform structure. Therefore, when forming a steel sheet, strain locally concentrates at the grain boundaries of the ferrite phase and the martensite phase, making it difficult to achieve excellent bendability and stretch flangeability. Therefore, the amount of temperature drop from the cooling stop temperature of the steel sheet is in the range of 0 to 100 ° C.
鋼板滞留時間：100秒以上1000秒以下
冷却停止後の滞留時間は、低温変態相の生成の進行にかかわり、ベイナイト相の体積分率を制御する上で重要である。滞留時間が100秒に満たない場合、滞留中のベイナイト変態が不十分となり、滞留後の室温までの冷却過程において、未変態オーステナイト相がマルテンサイト相となり、マルテンサイト相の体積分率が過剰となる結果、鋼板が過度に高強度化して延性が低下する。一方、滞留時間が1000秒を超えると、ベイナイト相の体積分率が過剰となり、延性および曲げ性が共に劣化する。1180MPa級以上の引張強度を確保すると共に優れた曲げ性、伸びフランジ性および延性を達成するには、滞留時間は100秒以上1000秒以下の範囲とする必要がある。
Steel plate residence time: 100 seconds or more and 1000 seconds or less The residence time after the cooling stop is related to the progress of the generation of the low temperature transformation phase, and is important for controlling the volume fraction of the bainite phase. When the residence time is less than 100 seconds, the bainite transformation during residence is insufficient, and in the cooling process to room temperature after residence, the untransformed austenite phase becomes the martensite phase, and the volume fraction of the martensite phase is excessive. As a result, the steel sheet becomes excessively strong and the ductility decreases. On the other hand, when the residence time exceeds 1000 seconds, the volume fraction of the bainite phase becomes excessive and both ductility and bendability deteriorate. In order to secure a tensile strength of 1180 MPa class or higher and achieve excellent bendability, stretch flangeability and ductility, the residence time needs to be in the range of 100 seconds to 1000 seconds.
冷却停止後の鋼板を上記滞留温度に保持する手段としては、例えば、焼鈍後の冷却設備の下工程に保温装置等を設けて、鋼板の温度を上記滞留温度に調整する手段等が挙げられる。ただし、冷却停止後の鋼板を上記滞留時間放冷しても冷却停止温度からの鋼板温度降下量が０〜100℃の温度範囲に収まる場合には、上記保温装置等の特別な手段を設けなくてもよい。なお、滞留後の鋼板は、従来公知の任意の方法により所望の温度に冷却される。 Examples of the means for holding the steel plate after stopping the cooling at the residence temperature include a means for adjusting the temperature of the steel plate to the residence temperature by providing a heat retaining device or the like in the lower process of the cooling equipment after annealing. However, if the steel plate temperature drop from the cooling stop temperature falls within the temperature range of 0 to 100 ° C even if the steel plate after the cooling stop is allowed to cool, the special means such as the heat retaining device is not provided. May be. In addition, the steel plate after a residence is cooled to desired temperature by the conventionally well-known arbitrary methods.
上記の焼鈍後、最終的に得られた冷延鋼板に、形状矯正や表面粗度調整の目的から調質圧延（スキンパス圧延）を行ってもかまわないが、過度にスキンパス圧延をすると鋼板に歪が導入されるため、結晶粒が展伸されて圧延加工組織となり、延性が低下するおそれがある。そのため、スキンパス圧延の圧下率は0.05％以上0.5％以下程度とすることが好ましい。 After the annealing described above, the cold-rolled steel sheet finally obtained may be subjected to temper rolling (skin pass rolling) for the purpose of shape correction or surface roughness adjustment. Therefore, the crystal grains are expanded to form a rolled structure, and the ductility may be reduced. Therefore, the rolling reduction of skin pass rolling is preferably about 0.05% to 0.5%.
（実施例１）
表１に示す成分組成になる鋼を溶製してスラブとし、1250℃に加熱後、表２に示す条件で、熱間仕上げ圧延を施し、塩酸酸洗後、冷間圧延、焼鈍処理および制御冷却処理を行い、板厚：1.0〜2.0mmの冷延鋼板を製造した。得られた冷延鋼板について、下記に示す材料試験により材料特性を調査した。得られた結果を表３に示す。
Example 1
Steels with the composition shown in Table 1 are melted to form slabs, heated to 1250 ° C, hot-finished under the conditions shown in Table 2, pickled with hydrochloric acid, cold-rolled, annealed and controlled. Cooling treatment was performed to produce a cold-rolled steel sheet having a thickness of 1.0 to 2.0 mm. About the obtained cold-rolled steel plate, the material characteristic was investigated by the material test shown below. The obtained results are shown in Table 3.
（１）鋼板の組織は、圧延方向の断面で、板厚の1/4程度の厚み付近の面を光学顕微鏡または走査型電子顕微鏡（SEM）で観察することにより調査した。観察はN=５（観察視野５箇所）で実施した。フェライト相の結晶粒径は、JIS G 0552（1998）に規定の方法（切断法）に準拠して結晶粒度を求め、平均粒径に換算した。ベイナイト相、マルテンサイト相についても同様に測定した。フェライト相、ベイナイト相およびマルテンサイト相の体積分率は倍率1000倍の断面組織写真を用いて、任意に設定した100mm×100mm四方の正方形領域内に存在する各相の占有面積を画像解析により求め、これを各相の体積分率とした。ベイナイト相とマルテンサイト相の区別は倍率3000倍の断面組織写真を用いて、フェライト相以外の低温変態相において炭化物の観察されるものをベイナイト相、炭化物の観察されない平滑な表面として観察されるものをマルテンサイト相、さらにセメンタイト相が層状に観察された場合パーライト相と判定した。なお、残留オーステナイト相は、Ｘ線回折法でMoのKα線を用いて求めた。また、残留オーステナイト相の体積分率は、上記と同様に圧延方向の断面で、板厚の1/4程度の付近の面を測定面とする試験片を用い、オーステナイト相の（２１１）面および（２２０）面と、フェライト相の（２００）面および（２２０）面とのピーク強度を測定して算出した。 (1) The structure of the steel sheet was examined by observing a surface near a thickness of about 1/4 of the sheet thickness with an optical microscope or a scanning electron microscope (SEM) in a cross section in the rolling direction. Observation was carried out at N = 5 (5 observation fields). The crystal grain size of the ferrite phase was calculated according to the method (cutting method) defined in JIS G 0552 (1998) and converted to an average grain size. It measured similarly about the bainite phase and the martensite phase. The volume fraction of the ferrite phase, bainite phase, and martensite phase is determined by image analysis to determine the area occupied by each phase existing in a 100 mm x 100 mm square area set arbitrarily using a cross-sectional structure photograph at a magnification of 1000 times. This was defined as the volume fraction of each phase. The distinction between the bainite phase and the martensite phase is made using a cross-sectional structure photograph at a magnification of 3000 times. When a martensite phase and further a cementite phase were observed in a layered state, it was determined to be a pearlite phase. The residual austenite phase was determined using Mo Kα rays by X-ray diffraction. Further, the volume fraction of the retained austenite phase is a cross section in the rolling direction in the same manner as described above, and a test piece whose surface is about 1/4 of the plate thickness is used as a measurement surface. The peak intensity between the (220) plane and the (200) plane and (220) plane of the ferrite phase was measured and calculated.
（２）引張特性は、圧延方向と90°の方向を長手方向（引張方向）とするJIS Z 2201に記載の５号試験片を用い、JIS Z 2241に準拠した引張試験を行い評価した。なお、引張特性の評価基準はTS×El≧17000MPa・％以上（TS：引張強度（MPa）、El：全伸び（％））を良好とした。 (2) Tensile properties were evaluated by performing a tensile test in accordance with JIS Z 2241 using No. 5 test piece described in JIS Z 2201 with the rolling direction and 90 ° as the longitudinal direction (tensile direction). The evaluation criteria for tensile properties were TS × El ≧ 17000 MPa ·% or more (TS: tensile strength (MPa), El: total elongation (%)).
（３）伸びフランジ性（穴拡げ率）は、日本鉄鋼連盟規格JFS T 1001に基づいて実施した。初期直径d0=10mmの穴を打抜き、60°の円錐ポンチを上昇させて穴を拡げた際に、亀裂が板厚を貫通したところでポンチ上昇を止め、亀裂貫通後の打抜き穴径dを測定し、穴拡げ率：（（d- d0）/ d0）×100（％）として算出した。同一番号の鋼板について３回試験を実施し、穴拡げ率の平均値（λ）を求めた。なお、伸びフランジ性（穴拡げ率）の評価基準はTS×λ≧25000MPa・％以上（TS：引張強度（MPa））を良好とした。 (3) Stretch flangeability (hole expansion rate) was carried out based on the Japan Iron and Steel Federation standard JFS T 1001. When a hole with an initial diameter of d 0 = 10 mm is punched and the hole is expanded by raising a 60 ° conical punch, the punch stops rising when the crack penetrates the plate thickness, and the punched hole diameter d is measured after the crack has penetrated. and, hole expansion ratio: ((d- d 0) / d 0) was calculated as × 100 (%). Three tests were performed on the same number of steel plates, and the average value (λ) of the hole expansion rate was obtained. The evaluation criteria for stretch flangeability (hole expansion ratio) was TS × λ ≧ 25000 MPa ·% or more (TS: tensile strength (MPa)).
（４）曲げ性は、JIS Z 2248に基づく押し曲げ法による曲げ試験にて評価した。
なお、板厚：1.0〜2.0mmの鋼板を用い、曲げ部の稜線と鋼板の圧延方向が平行になるようにサンプルを採取し、サンプルサイズは板厚×幅が40mm×150mm（サンプルの長手が圧延直角方向）とした。
押し金具の先端部の半径（ｒ）を種々変化させ、鋼板サンプルのわん曲部の外側の割れの有無を観察し、鋼板が割れることなく成形できる限界曲げ半径（Ｒ）と板厚（ｔ）の比である限界曲げ半径（Ｒ）／板厚（ｔ）：限界曲げ指数を求めた。
(4) The bendability was evaluated by a bending test by a push bending method based on JIS Z 2248.
In addition, using a steel plate with a thickness of 1.0 to 2.0 mm, take a sample so that the ridge line of the bent part and the rolling direction of the steel plate are parallel, and the sample size is 40 mm x 150 mm (thickness x width 40 mm x 150 mm). The direction perpendicular to rolling).
The bending radius (R) and thickness (t) at which the steel plate can be formed without cracking by changing the radius (r) of the front end of the metal fittings and observing the presence or absence of cracks outside the curved portion of the steel plate sample Ratio of critical bending radius (R) / plate thickness (t): The critical bending index was determined.
表３より、本発明例では、TS×El≧17000MPa・％以上、TS×λ≧25000MPa・％以上をいずれも満足するのみならず、限界曲げ指数が2.5以下という優れた曲げ性を有する高強度冷延鋼板が得られていることが分かる。
一方、鋼成分が本発明範囲外であるNo.５および６は所望の組織を発現することができず加工性に劣っていた。
また、以下の本発明範囲外の比較例についても、熱間仕上げ圧延温度が低いNo.７、熱間仕上げ圧延後の冷却速度が遅いNo.９、巻き取り温度が高いNo.10、冷間圧延率が低いNo.11、焼鈍後の冷却速度が遅いNo.15および冷却停止温度が低いNo.17は、いずれもフェライト相の体積分率が高く、引張強度が1180MPaを満足していない。
From Table 3, in the present invention example, TS × El ≧ 17000 MPa ·% or more and TS × λ ≧ 25000 MPa ·% or more are all satisfied, and the high bending strength with excellent bendability is 2.5 or less. It can be seen that a cold-rolled steel sheet is obtained.
On the other hand, Nos. 5 and 6 whose steel components were outside the scope of the present invention could not express the desired structure and were inferior in workability.
Moreover, also about the following comparative examples outside the scope of the present invention, No. 7 with a low hot finish rolling temperature, No. 9 with a slow cooling rate after hot finish rolling, No. 10 with a high winding temperature, cold No. 11 with a low rolling rate, No. 15 with a slow cooling rate after annealing, and No. 17 with a low cooling stop temperature all have a high volume fraction of the ferrite phase and do not satisfy the tensile strength of 1180 MPa.
さらに、熱間仕上げ圧延温度が高いNo.８、滞留中の温度降下量が多いNo.19および滞留時間の短いNo.20は、いずれもマルテンサイト相の体積分率が高く、加工性および曲げ性が共に劣っていた。 In addition, No. 8 with a high hot finish rolling temperature, No. 19 with a large temperature drop during residence, and No. 20 with a short residence time both have a high volume fraction of martensite phase, workability and bending. Both sexes were inferior.
冷間圧延率が高いNo.12、焼鈍温度が高いNo.14および焼鈍後の冷却速度が速いNo.16は、いずれもフェライト相の体積分率が少なく、加工性および曲げ性が共に劣っていた。 No. 12 with a high cold rolling rate, No. 14 with a high annealing temperature, and No. 16 with a fast cooling rate after annealing have a low volume fraction of ferrite phase, and both workability and bendability are inferior. It was.
焼鈍温度が低いNo.13は、ベイナイト相およびフェライト相の体積分率が本発明の範囲を外れかつ再結晶が不十分であるため、加工性に劣っていた。また、冷却停止温度が高いNo.18および滞留時間の長いNo.21は、ベイナイト相の体積分率が高く、加工性および曲げ性が共に劣っていた。 No. 13 having a low annealing temperature was inferior in workability because the volume fractions of the bainite phase and the ferrite phase were outside the range of the present invention and recrystallization was insufficient. Further, No. 18 having a high cooling stop temperature and No. 21 having a long residence time had a high volume fraction of the bainite phase, and both workability and bendability were inferior.
本発明に従い、鋼板中のＣ量を低減し、Cu、Ni、Cr、Mo、Ｖなど高価な元素を積極的に含有させずとも、ベイナイト相、フェライト相、マルテンサイト相、残留オーステナイト相の各々の体積分率を規定することにより、安価でかつ優れた曲げ性を有し、しかも引張強度（TS）が1180MPa以上の高強度冷延鋼板を得ることができる。また、本発明の高強度冷延鋼板は、特に厳しい形状に曲げ加工される自動車部品に好適であるが、自動車部品以外にも、建築および家電分野などの厳しい寸法精度や曲げ性が必要とされる用途にも好適である。 According to the present invention, each of the bainite phase, ferrite phase, martensite phase, and retained austenite phase can be achieved without reducing the amount of C in the steel sheet and positively containing expensive elements such as Cu, Ni, Cr, Mo, and V. By prescribing the volume fraction, it is possible to obtain a high-strength cold-rolled steel sheet that is inexpensive and has excellent bendability and has a tensile strength (TS) of 1180 MPa or more. The high-strength cold-rolled steel sheet of the present invention is particularly suitable for automobile parts that are bent into a strict shape. However, in addition to automobile parts, strict dimensional accuracy and bendability are required in the fields of architecture and home appliances. It is also suitable for use.
Claims (2)
Ｃ：0.10％以上0.15％以下、
Si：1.0％以上2.0％以下、
Mn：2.0％以上3.0％以下、
Ｐ：0.030％以下、
Ｓ：0.0050％以下、
Al：0.005％以上0.1％以下、
Ｎ：0.01％以下、
Ti：0.005％以上0.050％以下 および
Ｂ：0.0001％以上0.0050％以下
を含有し、残部がFe及び不可避的不純物からなる成分組成を有し、体積分率で、
ベイナイト相が50％以上70％以下、
フェライト相が20％以上40％以下、
マルテンサイト相が１％以上10％以下および
残留オーステナイト相が５％以下
を含む組織を有し、引張強度：1180MPa以上であることを特徴とする、曲げ性に優れた高強度冷延鋼板。 % By mass
C: 0.10% to 0.15%,
Si: 1.0% to 2.0%,
Mn: 2.0% to 3.0%,
P: 0.030% or less,
S: 0.0050% or less,
Al: 0.005% to 0.1%,
N: 0.01% or less,
Ti: 0.005% or more and 0.050% or less and B: 0.0001% or more and 0.0050% or less, with the balance being composed of Fe and inevitable impurities,
The bainite phase is 50% to 70%,
Ferrite phase is 20% to 40%,
A high-strength cold-rolled steel sheet having excellent bendability, having a structure including a martensite phase of 1% to 10% and a residual austenite phase of 5% or less and a tensile strength of 1180 MPa or more.
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