JP2010065272A - High-strength steel sheet and method for manufacturing the same - Google Patents
High-strength steel sheet and method for manufacturing the same Download PDFInfo
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本発明は、自動車、電気等の産業分野で使用される加工性に優れた引張強さ（ＴＳ）が９８０ＭＰａ以上の高強度鋼板およびその製造方法に関するものである。 The present invention relates to a high-strength steel sheet having excellent workability and having a tensile strength (TS) of 980 MPa or more used in industrial fields such as automobiles and electricity, and a method for producing the same.
近年、地球環境保全の見地から、自動車の燃費向上が重要な課題となっている。このため、車体材料の高強度化により車体部品の薄肉化を図り、車体そのものを軽量化しようとする動きが活発である。 In recent years, improving the fuel efficiency of automobiles has become an important issue from the viewpoint of global environmental conservation. For this reason, efforts are being made to reduce the thickness of the vehicle body parts by increasing the strength of the vehicle body material and to reduce the weight of the vehicle body itself.
一般に、鋼板の高強度化を図るためには、鋼板の組織全体に対してマルテンサイトやベイナイトなどの硬質相の割合を増加させる必要がある。しかしながら、硬質相の割合を増加させることによる鋼板の高強度化は加工性の低下を招くことから、高強度と優れた加工性を併せ持つ鋼板の開発が望まれている。これまでに、フェライト−マルテンサイト二相鋼（ＤＰ鋼）や残留オーステナイトの変態誘起塑性を利用したＴＲＩＰ鋼など、種々の複合組織鋼板が開発されてきた。 Generally, in order to increase the strength of a steel sheet, it is necessary to increase the ratio of hard phases such as martensite and bainite to the entire structure of the steel sheet. However, since increasing the strength of the steel sheet by increasing the proportion of the hard phase causes a decrease in workability, development of a steel sheet having both high strength and excellent workability is desired. To date, various composite steel sheets such as ferrite-martensitic duplex steel (DP steel) and TRIP steel utilizing transformation-induced plasticity of retained austenite have been developed.
複合組織鋼板において硬質相の割合を増加させた場合、鋼板の加工性は硬質相の加工性の影響を強く受けるようになる。これは、硬質相の割合が少なく軟質なポリゴナルフェライトが多い場合には、ポリゴナルフェライトの変形能が鋼板の加工性に対して支配的であり、硬質相の加工性が十分でない場合においても延性等の加工性は確保されたが、硬質相の割合が多い場合には、ポリゴナルフェライトの変形ではなく硬質相の変形能自体が鋼板の成形性に直接影響するようになり、硬質相自体の加工性が十分でないと、鋼板の加工性の劣化が著しくなるためである。 When the ratio of the hard phase is increased in the composite structure steel plate, the workability of the steel plate is strongly influenced by the workability of the hard phase. This is because when the ratio of hard phase is small and soft polygonal ferrite is large, the deformability of polygonal ferrite dominates the workability of the steel sheet, and even when the hard phase has insufficient workability. Workability such as ductility was ensured, but when the ratio of the hard phase is large, the deformability of the hard phase itself directly affects the formability of the steel sheet, not the deformation of polygonal ferrite, and the hard phase itself This is because if the workability is insufficient, the workability of the steel sheet is significantly deteriorated.
このため、冷延鋼板の場合には、焼鈍およびその後の冷却過程で生成するポリゴナルフェライトの量を調整する熱処理を行った後、鋼板を水焼入れしてマルテンサイトを生成させ、再び鋼板を昇温して高温保持することにより、マルテンサイトを焼戻して、硬質相であるマルテンサイト中に炭化物を生成させて、マルテンサイトの加工性を向上させてきた。しかしながら、このようなマルテンサイトの焼入れ・焼戻しには、例えば、水焼入れ機能を有する連続焼鈍設備のような特別な製造設備が必要となる。従って、鋼板を水焼入れした後、再び昇温して高温保持することができない通常の製造設備を使用する場合には、鋼板を高強度化することはできるが、硬質相であるマルテンサイトの加工性を向上させることはできなかった。 For this reason, in the case of a cold-rolled steel sheet, after performing annealing and adjusting the amount of polygonal ferrite generated in the subsequent cooling process, the steel sheet is water-quenched to generate martensite, and the steel sheet is raised again. By heating and holding at a high temperature, martensite has been tempered to generate carbides in the martensite that is a hard phase, thereby improving the workability of martensite. However, such martensite quenching and tempering requires special production equipment such as continuous annealing equipment having a water quenching function. Therefore, after water quenching the steel sheet, when using normal production equipment that cannot be heated again and kept at a high temperature, the steel sheet can be strengthened, but the martensite that is a hard phase can be processed. The property could not be improved.
また、マルテンサイト以外を硬質相とする鋼板として、主相をポリゴナルフェライト、硬質相をベイナイトやパーライトとし、かつこれらの硬質相であるベイナイトやパーライトに炭化物を生成させた鋼板がある。この鋼板は、ポリゴナルフェライトのみで加工性を向上させるのではなく、硬質相中に炭化物を生成させることにより硬質相自体の加工性も向上させ、特に、伸びフランジ性の向上を図る鋼板である。しかしながら、主相をポリゴナルフェライトとしている以上、引張強さ（ＴＳ）で９８０ＭＰａ以上の高強度化と加工性の両立を図ることは困難である。また、硬質相中に炭化物を生成させることによって硬質相自体の加工性の向上させても、ポリゴナルフェライトの加工性の良さには劣るため、引張強さ（ＴＳ）で９８０ＭＰａ以上の高強度化を図るためにポリゴナルフェライトの量を低減すると、十分な加工性を得ることができなくなる。 Further, as a steel plate having a hard phase other than martensite, there is a steel plate in which the main phase is polygonal ferrite, the hard phase is bainite or pearlite, and carbides are generated in these hard phases bainite or pearlite. This steel sheet is a steel sheet that not only improves the workability with polygonal ferrite alone, but also improves the workability of the hard phase itself by generating carbides in the hard phase, and in particular, improves the stretch flangeability. . However, as long as the main phase is polygonal ferrite, it is difficult to achieve both high strength and workability of 980 MPa or higher in tensile strength (TS). Also, even if the hard phase itself is improved in workability by generating carbides in the hard phase, the processability of polygonal ferrite is inferior, so the tensile strength (TS) is higher than 980 MPa. If the amount of polygonal ferrite is reduced in order to achieve the above, sufficient workability cannot be obtained.
特許文献１には、合金成分を規定し、鋼組織を、残留オーステナイトを有する微細で均一なベイナイトとすることにより、曲げ加工性および衝撃特性に優れる高張力鋼板が提案されている。
特許文献２には、所定の合金成分を規定し、鋼組織を、残留オーステナイトを有するベイナイトとし、かつベイナイト中の残留オーステナイト量を規定することにより、焼付け硬化性に優れた複合組織鋼板が提案されている。
特許文献３には、所定の合金成分を規定し、鋼組織を、残留オーステナイトを有するベイナイトを面積率で９０％以上、ベイナイト中の残留オーステナイト量を１％以上１５％以下とし、かつベイナイトの硬度(ＨＶ)を規定することにより、耐衝撃性に優れた複合組織鋼板が提案されている。
しかしながら、上述した鋼板には以下に述べる課題がある。
特許文献１に記載される成分組成では、鋼板に歪みを付与した際に、高歪域でのＴＲＩＰ効果を発現する安定した残留オーステナイトの量を確保することが困難であり、曲げ性は得られるものの、塑性不安定が生じるまでの延性が低く、張り出し性に劣る。
However, the above-described steel sheet has the following problems.
In the component composition described in Patent Document 1, it is difficult to ensure a stable amount of retained austenite that exhibits the TRIP effect in a high strain region when strain is applied to the steel sheet, and bendability is obtained. However, the ductility until plastic instability occurs is low, and the stretchability is inferior.
特許文献２に記載の鋼板は、焼付硬化性は得られるものの引張強さ（ＴＳ）を９８０ＭＰａ以上あるいはさらに１０５０ＭＰａ以上に高強度化しようとしても、ベイナイトあるいはさらにフェライトを主体とするマルテンサイトを極力抑制した組織であるため、強度の確保あるいは高強度化時における延性などの加工性の確保が困難である。 Although the steel sheet described in Patent Document 2 has bake hardenability, even when trying to increase the tensile strength (TS) to 980 MPa or more, or even 1050 MPa or more, bainite or martensite mainly composed of ferrite is suppressed as much as possible. Therefore, it is difficult to ensure the workability such as ductility when securing the strength or increasing the strength.
特許文献３に記載の鋼板は、耐衝撃性を向上させることを主目的としており、硬さがＨＶ ２５０以下のベイナイトを主相とし、具体的にはこれを９０％超で含む組織であるため、引張強さ（ＴＳ）を９８０ＭＰａ以上とすることは難しい。 The steel sheet described in Patent Document 3 is mainly intended to improve impact resistance, and has a main phase of bainite having a hardness of HV 250 or less, specifically a structure containing more than 90%. It is difficult to set the tensile strength (TS) to 980 MPa or more.
本発明は、上記の課題を有利に解決するもので、加工性に優れる引張強さ（ＴＳ）が９８０ＭＰａ以上の高強度鋼板を、その有利な製造方法とともに提供することを目的とする。
本発明の高強度鋼板には、鋼板の表面に溶融亜鉛めっきまたは合金化溶融亜鉛めっきを施した鋼板を含むものとする。
なお、本発明において、加工性に優れるとは、ＴＳ×Ｔ．ＥＬの値が２００００ＭＰａ・％以上、より好ましくは、２５０００ＭＰａ・％以上を満足することであるものとする。
An object of the present invention is to solve the above-mentioned problem advantageously, and an object of the present invention is to provide a high-strength steel sheet having a tensile strength (TS) excellent in workability of 980 MPa or more together with its advantageous production method.
The high-strength steel sheet of the present invention includes a steel sheet obtained by subjecting the surface of the steel sheet to hot dip galvanization or galvannealing.
In the present invention, excellent workability means that TS × T. It is assumed that the EL value satisfies 20000 MPa ·% or more, more preferably 25000 MPa ·% or more.
発明者らは、上記の課題を解決すべく、鋼板の成分組成およびミクロ組織について鋭意検討を重ねた。その結果、マルテンサイト組織を活用して高強度化を図るとともに、鋼板中のＣ量を０．１６％以上とＣ含有量を多くした上で、上部ベイナイト変態を活用することで、ＴＲＩＰ効果を得る上で有利な安定した残留オーステナイトを確保でき、優れた成形性を有する引張強さが９８０ＭＰａ以上の高強度鋼板が得られることを見出した。 In order to solve the above-mentioned problems, the inventors have conducted intensive studies on the component composition and microstructure of the steel sheet. As a result, the martensite structure is utilized to increase the strength, the C content in the steel sheet is increased to 0.16% or more and the C content is increased, and the upper bainite transformation is utilized to achieve the TRIP effect. It has been found that stable retained austenite advantageous in obtaining can be secured, and a high-strength steel sheet having excellent formability and a tensile strength of 980 MPa or more can be obtained.
本発明は、上記の知見に立脚するものであり、その要旨構成は次のとおりである。
１．質量％で
Ｃ：０.１６％以上０.７２％以下、
Ｓｉ：３.０％以下、
Ｍｎ：０.５％以上３.０％以下、
Ｐ：０．１％以下、
Ｓ：０．０７％以下、
Ａｌ：３.０％以下および
Ｎ：０．０１０％以下
を含有し、かつＳｉ＋Ａｌが０.７％以上を満足し、残部はＦｅおよび不可避不純物の組成になり、
鋼板組織として、マルテンサイトの鋼板組織全体に対する面積率が１０％以上９０％以下、残留オーステナイト量が５％以上５０％以下、上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率が５％以上であり、前記マルテンサイトの鋼板組織全体に対する面積率、前記残留オーステナイト量および前記上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率の合計が６０％以上、ポリゴナルフェライトの鋼板組織全体に対する面積率が１０％以下（０％を含む）を満足し、かつ前記残留オーステナイト中の平均Ｃ量が０.７０％以上２.００％以下であり、引張強さが９８０ＭＰａ以上であることを特徴とする高強度鋼板。
The present invention is based on the above findings, and the gist of the present invention is as follows.
1. C: 0.16% to 0.72% by mass%,
Si: 3.0% or less,
Mn: 0.5% to 3.0%,
P: 0.1% or less,
S: 0.07% or less,
Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, the balance is the composition of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of martensite to the entire steel sheet structure is 10% to 90%, the amount of retained austenite is 5% to 50%, and the area ratio of bainitic ferrite in the upper bainite to the entire steel sheet structure is 5%. The total area ratio of the martensite to the entire steel sheet structure, the amount of retained austenite, and the area ratio of the bainitic ferrite in the upper bainite to the entire steel sheet structure is 60% or more, and the entire steel structure of polygonal ferrite That the area ratio with respect to the above is 10% or less (including 0%), the average C content in the retained austenite is 0.70% or more and 2.00% or less, and the tensile strength is 980 MPa or more. High-strength steel sheet characterized.
２．前記鋼板がさらに、質量％で、
Ｃｒ：０.０５％以上５.０％以下、
Ｖ：０．００５％以上１．０％以下および
Ｍｏ：０．００５％以上０．５％以下
のうちから選ばれる１種または２種以上の元素を含有することを特徴とする上記１に記載の高強度鋼板。
2. The steel sheet is further in mass%,
Cr: 0.05% to 5.0%,
2. One or more elements selected from V: 0.005% to 1.0% and Mo: 0.005% to 0.5% are described in 1 above High strength steel plate.
３．前記鋼板がさらに、質量％で、
Ｔｉ：０．０１％以上０．１％以下および
Ｎｂ：０．０１％以上０．１％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする上記１または２に記載の高強度鋼板。
3. The steel sheet is further in mass%,
1 or 2 above, which contains one or two elements selected from Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%. High strength steel sheet as described.
４．前記鋼板がさらに、質量％で、
Ｂ：０．０００３％以上０．００５０％以下
を含有することを特徴とする上記１乃至３のいずれか１項に記載の高強度鋼板。
4). The steel sheet is further in mass%,
B: The high-strength steel sheet according to any one of 1 to 3 above, which contains 0.0003% or more and 0.0050% or less.
５．前記鋼板がさらに、質量％で、
Ｎｉ：０．０５％以上２.０％以下および
Ｃｕ：０．０５％以上２.０％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする上記１乃至４のいずれか１項に記載の高強度鋼板。
5). The steel sheet is further in mass%,
1 to 4 above, which contains one or two elements selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less. The high-strength steel sheet according to any one of the items.
６．前記鋼板がさらに、質量％で、
Ｃａ：０．００１％以上０．００５％以下および
ＲＥＭ：０．００１％以上０．００５％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする上記１乃至５のいずれか１項に記載の高強度鋼板。
6). The steel sheet is further in mass%,
1 to 5 above, which contains one or two elements selected from Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%. The high-strength steel sheet according to any one of the items.
７．上記１乃至６のいずれか１項に記載の鋼板の表面に、溶融亜鉛めっき層または合金化溶融亜鉛めっき層を具えることを特徴とする高強度鋼板。 7). 7. A high-strength steel sheet comprising a hot-dip galvanized layer or an alloyed hot-dip galvanized layer on the surface of the steel sheet according to any one of 1 to 6 above.
８．上記１乃至６のいずれか１項に記載の成分組成になる鋼片を、熱間圧延し、冷間圧延により冷延鋼板とし、ついで該冷延鋼板を、オーステナイト単相域で１５秒以上６００秒以下焼鈍した後、３５０℃以上４９０℃以下の保持温度域で定める冷却停止温度：Ｔ℃まで冷却するに際し、少なくとも５５０℃までは平均冷却速度を５℃/ｓ以上に制御して冷却し、該保持温度域で１５秒以上１０００秒以下保持することを特徴とする高強度鋼板の製造方法。 8). The steel slab having the composition described in any one of 1 to 6 above is hot-rolled to form a cold-rolled steel sheet by cold rolling, and then the cold-rolled steel sheet is 600 seconds or more in the austenite single-phase region. After annealing for a second or less, a cooling stop temperature determined in a holding temperature range of 350 ° C. or more and 490 ° C. or less: When cooling to T ° C., at least 550 ° C. is controlled by controlling the average cooling rate to 5 ° C./s or more, and A method for producing a high-strength steel sheet, characterized by holding for 15 seconds to 1000 seconds in the holding temperature range.
９．前記冷却停止温度：Ｔ℃までの冷却時もしくは前記保持温度域で、溶融亜鉛めっき処理または合金化溶融亜鉛めっき処理を施すことを特徴とする上記８に記載の高強度鋼板の製造方法。 9. 9. The method for producing a high-strength steel sheet as described in 8 above, wherein a hot dip galvanizing treatment or an alloying hot dip galvanizing treatment is performed at the cooling stop temperature: T ° C. or in the holding temperature range.
本発明によれば、加工性に優れる引張強さ（ＴＳ）が９８０ＭＰａ以上の高強度鋼板をその有利な製造方法とともに提供することができ、自動車、電気等の産業分野での利用価値は非常に大きく、特に自動車車体の軽量化に対して極めて有用である。 According to the present invention, it is possible to provide a high-strength steel sheet having an excellent workability and a tensile strength (TS) of 980 MPa or more together with its advantageous production method, and the utility value in the industrial fields such as automobiles and electricity is very high. It is very useful especially for reducing the weight of automobile bodies.
以下、本発明を具体的に説明する。
まず、本発明において、鋼板組織を上記のように限定した理由について述べる。以下、面積率は、鋼板組織全体に対する面積率とする。
The present invention will be specifically described below.
First, the reason why the steel sheet structure is limited as described above in the present invention will be described. Hereinafter, the area ratio is the area ratio relative to the entire steel sheet structure.
マルテンサイトの面積率：１０％以上９０％以下
マルテンサイトは硬質相であり、鋼板を高強度化するために必要な組織である。マルテンサイトの面積率が１０％未満では、鋼板の引張強さ（ＴＳ）が９８０ＭＰａを満足しない。一方、マルテンサイトの面積率が９０％を超えると、上部ベイナイトが少なくなり、その結果、炭素が濃化した安定した残留オーステナイト量が確保できないため、延性等の加工性が低下することが問題となる。従って、マルテンサイトの面積率は、１０％以上９０％以下とする。なお、好ましくは８５％以下、さらに好ましくは７５％以下である。
Martensite area ratio: 10% or more and 90% or less Martensite is a hard phase and is a structure necessary for increasing the strength of a steel sheet. When the area ratio of martensite is less than 10%, the tensile strength (TS) of the steel sheet does not satisfy 980 MPa. On the other hand, when the area ratio of martensite exceeds 90%, the upper bainite is reduced, and as a result, a stable retained austenite amount in which carbon is concentrated cannot be secured, so that workability such as ductility is deteriorated. Become. Therefore, the area ratio of martensite is 10% or more and 90% or less. In addition, Preferably it is 85% or less, More preferably, it is 75% or less.
残留オーステナイト量：５％以上５０％以下
残留オーステナイトは、加工時にＴＲＩＰ効果によりマルテンサイト変態し、歪分散能を高めることにより延性を向上させる。
本発明の鋼板では、上部ベイナイト変態を活用して、特に、炭素濃化量を高めた残留オーステナイトを、上部ベイナイト中に形成せしめる。その結果、加工時に高歪域でもＴＲＩＰ効果を発現できる残留オーステナイトを得ることができる。このような残留オーステナイトとマルテンサイトを併存させて活用することにより、引張強さ（ＴＳ）が９８０ＭＰａ以上の高強度領域でも良好な加工性が得られ、具体的には、ＴＳ×Ｔ．Ｅｌの値を２００００ＭＰａ・％以上とすることができ、強度と延性のバランスに優れた鋼板を得ることができる。
ここで、上部ベイナイト中の残留オーステナイトは、上部ベイナイト中のベイニティックフェライトのラス間に形成され、細かく分布するため、組織観察によりその量（面積率）を求めるには高倍率で大量の測定が必要であり正確に定量することは難しい。しかし、該ベイニティックフェライトのラス間に形成される残留オーステナイトの量は、形成されるベイニティックフェライト量にある程度見合った量である。そこで、発明者らが検討した結果、上部ベイナイト中のベイニティックフェライトの面積率が５％以上で、かつ従来から行われている残留オーステナイト量を測定する手法であるＸ線回折（ＸＲＤ）による強度測定、具体的にはフェライトとオーステナイトのＸ線回折の強度比から求められる鋼板中の残留オーステナイト量が５％以上であれば、十分なＴＲＩＰ効果を得ることができ、引張強さ（ＴＳ）が９８０ＭＰａ以上で、ＴＳ×Ｔ．Ｅｌが２００００ＭＰａ・％以上を達成できることが分かった。なお、従来から行われている残留オーステナイト量の測定手法で得られた残留オーステナイト量は、残留オーステナイトの鋼板組織全体に対する面積率と同等であることを確認している。
残留オーステナイト量が５％未満の場合、十分なＴＲＩＰ効果が得られない。一方、５０％を超えると、ＴＲＩＰ効果発現後に生じる硬質なマルテンサイトが過大となり、靭性の劣化などが問題となる。従って、残留オーステナイトの量は、５％以上５０％以下の範囲とする。好ましくは５％超、より好ましくは１０％以上４５％以下の範囲である。さらに好ましくは、１５％以上４０％以下の範囲である。
Residual austenite amount: 5% or more and 50% or less Residual austenite undergoes martensitic transformation by the TRIP effect during processing, and improves ductility by increasing strain dispersibility.
In the steel sheet of the present invention, utilizing the upper bainite transformation, in particular, retained austenite having an increased carbon concentration is formed in the upper bainite. As a result, retained austenite that can exhibit the TRIP effect even in a high strain region during processing can be obtained. By utilizing such retained austenite and martensite in combination, good workability can be obtained even in a high strength region where the tensile strength (TS) is 980 MPa or more, specifically, TS × T. The value of El can be set to 20000 MPa ·% or more, and a steel sheet having an excellent balance between strength and ductility can be obtained.
Here, the retained austenite in the upper bainite is formed between the laths of the bainitic ferrite in the upper bainite and is finely distributed. Is necessary and accurate quantification is difficult. However, the amount of retained austenite formed between the laths of the bainitic ferrite is a certain amount commensurate with the amount of bainitic ferrite formed. Thus, as a result of investigations by the inventors, the area ratio of bainitic ferrite in the upper bainite is 5% or more, and X-ray diffraction (XRD) is a method for measuring the amount of retained austenite that has been conventionally performed. If the amount of retained austenite in the steel sheet obtained from strength measurement, specifically, the X-ray diffraction intensity ratio of ferrite and austenite is 5% or more, sufficient TRIP effect can be obtained, and tensile strength (TS) Of TS × T. It was found that El can achieve 20000 MPa ·% or more. It has been confirmed that the amount of retained austenite obtained by a conventional method for measuring the amount of retained austenite is equivalent to the area ratio of retained austenite to the entire steel sheet structure.
When the amount of retained austenite is less than 5%, a sufficient TRIP effect cannot be obtained. On the other hand, if it exceeds 50%, hard martensite generated after the TRIP effect appears becomes excessive, which causes problems such as deterioration of toughness. Therefore, the amount of retained austenite is in the range of 5% to 50%. Preferably it is more than 5%, more preferably in the range of 10% to 45%. More preferably, it is the range of 15% or more and 40% or less.
残留オーステナイト中の平均Ｃ量：０.７０％以上２.００％以下
ＴＲＩＰ効果を活用して優れた加工性を得るためには、引張強さ（ＴＳ）が９８０ＭＰａ〜２.５ＧＰａ級の高強度鋼板においては、残留オーステナイト中のＣ量が重要である。本発明の鋼板では、上部ベイナイト中のベイニティックフェライトのラス間に形成される残留オーステナイトにＣを濃化させる。該ラス間の残留オーステナイト中に濃化されるＣ量を正確に評価することは困難であるが、発明者らが検討した結果、本発明の鋼板においては、従来行われている残留オーステナイト中の平均Ｃ量（残留オーステナイト中のＣ量の平均）を測定する方法であるＸ線回折（ＸＲＤ）での回折ピークのシフト量から求める残留オーステナイト中の平均Ｃ量が０．７０％以上であれば、優れた加工性が得られることが分かった。
残留オーステナイト中の平均Ｃ量が０．７０％未満の場合、加工時において低歪域でマルテンサイト変態が生じてしまい、加工性を向上させる高歪域でのＴＲＩＰ効果が得られない。一方、残留オーステナイト中の平均Ｃ量が２．００％を超えると、残留オーステナイトが過剰に安定となり、加工中にマルテンサイト変態が生じず、ＴＲＩＰ効果が発現しないことにより、延性が低下する。従って、残留オーステナイト中の平均Ｃ量は０.７０％以上２.００％以下の範囲とする。好ましくは、０.９０％以上１.５０％以下の範囲である。
Average C content in retained austenite: 0.70% or more and 2.00% or less To obtain excellent workability by utilizing the TRIP effect, high tensile strength (TS) of 980 MPa to 2.5 GPa class In the steel sheet, the amount of C in the retained austenite is important. In the steel sheet of the present invention, C is concentrated in the retained austenite formed between the laths of bainitic ferrite in the upper bainite. Although it is difficult to accurately evaluate the amount of C concentrated in the retained austenite between the laths, as a result of the study by the inventors, in the steel sheet of the present invention, the conventional austenite in the retained austenite If the average C content in the retained austenite obtained from the shift amount of the diffraction peak in X-ray diffraction (XRD), which is a method for measuring the average C content (average of the C content in the retained austenite) is 0.70% or more It was found that excellent processability can be obtained.
When the average C content in the retained austenite is less than 0.70%, martensitic transformation occurs in the low strain region during processing, and the TRIP effect in the high strain region that improves workability cannot be obtained. On the other hand, if the average C content in the retained austenite exceeds 2.00%, the retained austenite becomes excessively stable, the martensitic transformation does not occur during processing, and the TRIP effect does not appear, thereby reducing ductility. Therefore, the average amount of C in the retained austenite is set in the range of 0.70% or more and 2.00% or less. Preferably, it is 0.90% or more and 1.50% or less of range.
上部ベイナイト中のベイニティックフェライトの面積率：５％以上
上部ベイナイト変態によるベイニティックフェライトの生成は、未変態オーステナイト中のＣを濃化させ、加工時に高歪域でＴＲＩＰ効果を発現して歪分解能を高める残留オーステナイトを得るために必要である。オーステナイトからベイナイトへの変態は、およそ１５０〜５５０℃の広い温度範囲にわたって起こり、この温度範囲内で生成するベイナイトには種々のものが存在する。従来技術では、このような種々のベイナイトを単にベイナイトと規定する場合が多かったが、本発明で目標とする加工性を得るためには、ベイナイト組織を明確に規定する必要があることから、上部ベイナイトおよび下部ベイナイトを次のように定義する。
上部ベイナイトは、ラス状のベイニティックフェライトと、ベイニッティクフェライトの間に存在する残留オーステナイトおよび／または炭化物とからなり、ラス状のベイニティックフェライト中に規則正しく並んだ細かな炭化物が存在しないことが特徴である。一方、下部ベイナイトは、ラス状のベイニティックフェライトと、ベイニッティクフェライトの間に存在する残留オーステナイトおよび／または炭化物とからなることは、上部ベイナイトと共通であるが、下部ベイナイトでは、ラス状のベイニティックフェライト中に規則正しく並んだ細かな炭化物が存在することが特徴である。
つまり、上部ベイナイトと下部ベイナイトは、ベイニティックフェライト中における規則正しく並んだ細かな炭化物の有無によって区別される。このようなベイニティックフェライト中における炭化物の生成状態の差は、残留オーステナイト中への炭素の濃化に大きな影響を与える。つまり、上部ベイナイトのベイニティックフェライトの面積率が５％未満の場合、ベイナイト変態を進めた場合においても、炭素はベイニティックフェライト中に炭化物として生成する量が多くなり、結果的にラス間に存在する残留オーステナイト中への炭素濃化量が減少して、加工時に高歪域でＴＲＩＰ効果を発現する残留オーステナイト量が減少することが問題となる。従って、上部ベイナイト中のベイニティックフェライトの面積率は、鋼板組織全体に対する面積率で５％以上必要である。一方、上部ベイナイト中のベイニティックフェライトの面積率が鋼板組織全体に対する面積率で８５％を超えると、強度の確保が困難となる場合があるため、８５％以下とすることが好ましい。
The area ratio of bainitic ferrite in the upper bainite: 5% or more The formation of bainitic ferrite by the upper bainite transformation concentrates C in the untransformed austenite and exhibits the TRIP effect in the high strain region during processing. It is necessary to obtain retained austenite that enhances strain resolution. The transformation from austenite to bainite occurs over a wide temperature range of approximately 150 to 550 ° C., and various types of bainite are produced within this temperature range. In the prior art, such various bainite was often simply defined as bainite, but in order to obtain the target workability in the present invention, it is necessary to clearly define the bainite structure. The bainite and lower bainite are defined as follows.
The upper bainite is composed of lath-like bainitic ferrite and residual austenite and / or carbide existing between bainitic ferrite, and there is no fine carbide regularly arranged in lath-like bainitic ferrite. It is a feature. On the other hand, the lower bainite is composed of the lath-shaped bainitic ferrite and the residual austenite and / or carbide existing between the bainitic ferrites in common with the upper bainite. It is characterized by the presence of fine carbides regularly arranged in the bainitic ferrite.
That is, the upper bainite and the lower bainite are distinguished by the presence or absence of regularly arranged fine carbides in bainitic ferrite. Such a difference in the formation state of carbides in bainitic ferrite has a great influence on the concentration of carbon in retained austenite. In other words, when the area ratio of the bainitic ferrite of the upper bainite is less than 5%, even when the bainite transformation is advanced, the amount of carbon produced as carbides in the bainitic ferrite increases, resulting in The amount of carbon enriched in the retained austenite present in the steel decreases, and the amount of retained austenite that exhibits the TRIP effect in the high strain region during processing decreases. Therefore, the area ratio of bainitic ferrite in the upper bainite needs to be 5% or more in terms of the area ratio with respect to the entire steel sheet structure. On the other hand, if the area ratio of bainitic ferrite in the upper bainite exceeds 85% in terms of the area ratio with respect to the entire steel sheet structure, it may be difficult to ensure the strength.
マルテンサイトの面積率、残留オーステナイト量および上部ベイナイト中のベイニティックフェライトの面積率の合計：６０％以上
マルテンサイトの面積率、残留オーステナイト量および上部ベイナイト中のベイニティックフェライトの面積率のそれぞれを上記した範囲で満足するだけでは不十分で、マルテンサイトの面積率、残留オーステナイト量および上部ベイナイト中のベイニティックフェライトの面積率の合計が６０％以上である必要がある。６０％未満の場合、強度不足や加工性の低下またはその両方を生じる。好ましくは７０％以上、より好ましくは８０％以上である。
Total ratio of martensite area ratio, residual austenite amount and area ratio of bainitic ferrite in upper bainite: 60% or more Each of martensite area ratio, residual austenite amount and area ratio of bainitic ferrite in upper bainite Is not sufficient to satisfy the above-mentioned range, and the sum of the area ratio of martensite, the amount of retained austenite and the area ratio of bainitic ferrite in the upper bainite needs to be 60% or more. If it is less than 60%, the strength is insufficient, the workability is lowered, or both. Preferably it is 70% or more, more preferably 80% or more.
ポリゴナルフェライト面積率：１０％以下（０％を含む）
ポリゴナルフェライトの面積率が１０％を超えると、引張強さ（ＴＳ）：９８０ＭＰａ以上を満足することが困難になると同時に、加工時に硬質組織内に混在した軟質なポリゴナルフェライトに歪が集中することにより加工時に容易に亀裂が発生し、結果として所望の加工性を得られない。ここで、ポリゴナルフェライトの面積率が１０％以下であれば、ポリゴナルフェライトが存在しても硬質相中に少量のポリゴナルフェライトが孤立分散した状態となり、歪の集中を抑制することができ、加工性の劣化を避けることができる。従って、ポリゴナルフェライトの面積率は１０％以下とする。好ましくは５％以下、さらに好ましくは３％以下であり、０％であってもよい。
Polygonal ferrite area ratio: 10% or less (including 0%)
When the area ratio of polygonal ferrite exceeds 10%, it becomes difficult to satisfy the tensile strength (TS) of 980 MPa or more, and at the same time, strain concentrates on the soft polygonal ferrite mixed in the hard structure during processing. As a result, cracks are easily generated during processing, and as a result, desired workability cannot be obtained. Here, if the area ratio of polygonal ferrite is 10% or less, even if polygonal ferrite is present, a small amount of polygonal ferrite is isolated and dispersed in the hard phase, and strain concentration can be suppressed. Degradation of workability can be avoided. Therefore, the area ratio of polygonal ferrite is 10% or less. Preferably it is 5% or less, More preferably, it is 3% or less, and 0% may be sufficient.
本発明の鋼板には、残部組織として、パーライトやウィドマンステッテンフェライト、下部ベイナイトを含んでも構わない。その場合、残部組織の許容含有量は、面積率で２０％以下とすることが好ましい。より好ましくは、１０％以下である。 The steel sheet of the present invention may contain pearlite, Widmanstatten ferrite, or lower bainite as the remaining structure. In that case, the allowable content of the remaining tissue is preferably 20% or less in terms of area ratio. More preferably, it is 10% or less.
次に、本発明において、鋼板の成分組成を上記のように限定した理由について述べる。なお、以下の成分組成を表す％は質量％を意味するものとする。
Ｃ：０.１６％以上０.７２％以下
Ｃは鋼板の高強度化および安定した残留オーステナイト量を確保するのに必要不可欠な元素であり、マルテンサイト量の確保および室温でオーステナイトを残留させるために必要な元素である。Ｃ量が、０．１６％未満では、鋼板の強度と加工性を確保することが難しい。一方、Ｃ量が０．７２％を超えると、溶接部および熱影響部の硬化が著しく溶接性が劣化する。従って、Ｃ量は０.１６％以上０.７２％以下の範囲とする。好ましくは、０．２０％を超え０.４８％以下の範囲であり、さらに好ましくは０．２５％以上である。
Next, the reason why the component composition of the steel sheet is limited as described above in the present invention will be described. In addition,% showing the following component compositions shall mean the mass%.
C: 0.16% or more and 0.72% or less C is an element indispensable for increasing the strength of a steel sheet and ensuring a stable amount of retained austenite, to ensure the amount of martensite and to retain austenite at room temperature. It is a necessary element. If the C content is less than 0.16%, it is difficult to ensure the strength and workability of the steel sheet. On the other hand, if the amount of C exceeds 0.72%, the welded portion and the heat affected zone are hardened and the weldability is deteriorated. Therefore, the C content is in the range of 0.16% or more and 0.72% or less. Preferably, it is 0.20% or more and 0.48% or less of range, More preferably, it is 0.25% or more.
Ｓｉ：３.０％以下（０％を含む）
Ｓｉは、固溶強化により鋼の強度向上に寄与する有用な元素である。しかしながら、Ｓｉ量が３．０％を超えると、ポリゴナルフェライトおよびベイニティックフェライト中への固溶量の増加による加工性、靭性の劣化を招き、また、赤スケール等の発生による表面性状の劣化や、溶融めっきを施す場合には、めっき付着性および密着性の劣化を引き起こす。従って、Ｓｉ量は３.０％以下とする。好ましくは２．６％以下である。さらに好ましくは、２．２％以下である。
また、Ｓｉは、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素であることから、Ｓｉ量は０．５％以上とすることが好ましいが、炭化物の生成をＡｌのみで抑制する場合には、Ｓｉは添加する必要はなく、Ｓｉ量は０％であっても良い。
Si: 3.0% or less (including 0%)
Si is a useful element that contributes to improving the strength of steel by solid solution strengthening. However, if the amount of Si exceeds 3.0%, the workability and toughness deteriorate due to the increase in the amount of solid solution in polygonal ferrite and bainitic ferrite, and the surface properties due to the occurrence of red scale, etc. In the case of deterioration or hot dipping, it causes deterioration of plating adhesion and adhesion. Therefore, the Si content is 3.0% or less. Preferably it is 2.6% or less. More preferably, it is 2.2% or less.
Si is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite. Therefore, the Si content is preferably 0.5% or more, but the formation of carbides is only Al. In the case of suppressing by Si, Si does not need to be added, and the Si amount may be 0%.
Ｍｎ：０.５％以上３.０％以下
Ｍｎは、鋼の強化に有効な元素である。Ｍｎ量が０．５％未満では、焼鈍後の冷却中にベイナイトやマルテンサイトが生成する温度よりも高い温度域で炭化物が析出するため、鋼の強化に寄与する硬質相の量を確保することができない。一方、Ｍｎ量が３.０％を超えると、鋳造性の劣化などを引き起こす。従って、Ｍｎ量は０.５％以上３.０％以下の範囲とする。好ましくは１.５％以上２.５％以下の範囲とする。
Mn: 0.5% to 3.0% Mn is an element effective for strengthening steel. If the amount of Mn is less than 0.5%, carbide precipitates in a temperature range higher than the temperature at which bainite and martensite are generated during cooling after annealing, so ensure the amount of hard phase that contributes to strengthening of steel. I can't. On the other hand, if the amount of Mn exceeds 3.0%, castability is deteriorated. Therefore, the amount of Mn is set in the range of 0.5% to 3.0%. Preferably it is 1.5 to 2.5% of range.
Ｐ：０．１％以下
Ｐは鋼の強化に有用な元素であるが、Ｐ量が０．１％を超えると、粒界偏析により脆化することにより耐衝撃性を劣化させ、鋼板に合金化溶融亜鉛めっきを施す場合には、合金化速度を大幅に遅延させる。従って、Ｐ量は０.１％以下とする。好ましくは０．０５％以下である。なお、Ｐ量は、低減することが好ましいが、０．００５％未満とするには大幅なコスト増加を引き起こすため、その下限は０．００５％程度とすることが好ましい。
P: 0.1% or less
P is an element useful for strengthening steel, but if the P content exceeds 0.1%, the impact resistance deteriorates due to embrittlement due to grain boundary segregation, and the steel sheet is subjected to alloying hot dip galvanizing. In this case, the alloying speed is greatly delayed. Therefore, the P content is 0.1% or less. Preferably it is 0.05% or less. The amount of P is preferably reduced, but if it is less than 0.005%, it causes a significant increase in cost, so the lower limit is preferably about 0.005%.
Ｓ：０．０７％以下
Ｓは、ＭｎＳを生成して介在物となり、耐衝撃性の劣化や溶接部のメタルフローに沿った割れの原因となるため、Ｓ量を極力低減することが好ましい。しかしながら、Ｓ量を過度に低減することは、製造コストの増加を招くため、Ｓ量は０．０７％以下とする。好ましくは０．０５％以下であり、より好ましくは０．０１％以下である。なお、Ｓを０．０００５％未満とするには大きな製造コストの増加を伴うため、製造コストの点からはその下限は０．０００５％程度である。
S: 0.07% or less
Since S generates MnS to become inclusions, causing deterioration of impact resistance and cracking along the metal flow of the weld, it is preferable to reduce the amount of S as much as possible. However, excessively reducing the amount of S causes an increase in manufacturing cost, so the amount of S is set to 0.07% or less. Preferably it is 0.05% or less, More preferably, it is 0.01% or less. In addition, since it is accompanied by a big increase in manufacturing cost to make S less than 0.0005%, the lower limit is about 0.0005% from the point of manufacturing cost.
Ａｌ：３.０％以下
Ａｌは、鋼の強化に有用な元素であるとともに、製鋼工程で脱酸剤として添加される有用な元素である。Ａｌ量が３.０％を超えると、鋼板中の介在物が多くなり延性を劣化させる。従って、Ａｌ量は３.０％以下とする。好ましくは、２．０％以下である。
また、Ａｌは、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素であり、また、脱酸効果を得るために、Ａｌ量は、０．００１％以上とすることが好ましく、より好ましくは０．００５％以上とする。なお、本発明におけるＡｌ量は、脱酸後に鋼板中に含有するＡｌ量とする。
Al: 3.0% or less Al is an element useful for strengthening steel and a useful element added as a deoxidizer in the steel making process. When the amount of Al exceeds 3.0%, the inclusions in the steel sheet increase and the ductility deteriorates. Therefore, the Al content is 3.0% or less. Preferably, it is 2.0% or less.
Al is an element useful for suppressing the formation of carbides and promoting the formation of retained austenite. In order to obtain a deoxidizing effect, the Al content should be 0.001% or more. Preferably, it is more preferably 0.005% or more. The amount of Al in the present invention is the amount of Al contained in the steel sheet after deoxidation.
Ｎ：０．０１０％以下
Ｎは、鋼の耐時効性を最も大きく劣化させる元素であり、極力低減することが好ましい。Ｎ量が０．０１０％を超えると耐時効性の劣化が顕著となるため、Ｎ量は０．０１０％以下とする。なお、Ｎを０．００１％未満とするには大きな製造コストの増加を招くため、製造コストの点からは、その下限は０．００１％程度である。
N: 0.010% or less
N is an element that most deteriorates the aging resistance of steel, and is preferably reduced as much as possible. When the N content exceeds 0.010%, deterioration of aging resistance becomes remarkable, so the N content is set to 0.010% or less. Note that, if N is less than 0.001%, a large increase in manufacturing cost is caused, so that the lower limit is about 0.001% from the viewpoint of manufacturing cost.
以上、基本成分について説明したが、本発明では、上記の成分範囲を満足するだけでは不十分で、次式を満足させる必要がある。
Ｓｉ＋Ａｌ≧０.７％
ＳｉおよびＡｌはともに、上記したように、炭化物の生成を抑制し、残留オーステナイトの生成を促進するのに有用な元素である。炭化物の生成の抑制は、ＳｉまたはＡｌを単独で含有させても効果があるが、Ｓｉ量とＡｌ量の合計で０．７％以上を満足する必要がある。なお、上掲式におけるＡｌ量は、脱酸後に鋼板中に含有するＡｌ量とする。
The basic component has been described above. However, in the present invention, it is not sufficient to satisfy the above component range, and it is necessary to satisfy the following equation.
Si + Al ≧ 0.7%
As described above, both Si and Al are useful elements for suppressing the formation of carbides and promoting the formation of retained austenite. Although suppression of the formation of carbides is effective even if Si or Al is contained alone, it is necessary to satisfy 0.7% or more in total of the Si amount and the Al amount. The amount of Al in the above formula is the amount of Al contained in the steel sheet after deoxidation.
また、本発明では上記した基本成分の他、以下に述べる成分を適宜含有させることができる。
Ｃｒ：０.０５％以上５.０％以下、Ｖ：０.００５％以上１.０％以下、Ｍｏ：０.００５％以上０.５％以下のうちから選ばれる１種または２種以上
Ｃｒ、ＶおよびＭｏは焼鈍温度からの冷却時にパーライトの生成を抑制する作用を有する元素である。その効果は、Ｃｒ：０.０５％以上、Ｖ：０.００５％以上およびＭｏ：０.００５％以上で得られる。一方、Ｃｒ：５.０％、Ｖ：１.０％およびＭｏ：０.５％を超えると、硬質なマルテンサイトの量が過大となり、必要以上に高強度となる。従って、Ｃｒ、ＶおよびＭｏを含有させる場合には、Ｃｒ：０．０５％以上５.０％以下、Ｖ：０．００５％以上１．０％以下、およびＭｏ：０．００５％以上０．５％以下の範囲とする。
Moreover, in this invention, the component described below other than the above-mentioned basic component can be contained appropriately.
Cr: 0.05% or more and 5.0% or less, V: 0.005% or more and 1.0% or less, Mo: one or more selected from 0.005% or more and 0.5% or less Cr , V and Mo are elements having an action of suppressing the formation of pearlite during cooling from the annealing temperature. The effect is obtained when Cr: 0.05% or more, V: 0.005% or more, and Mo: 0.005% or more. On the other hand, if it exceeds Cr: 5.0%, V: 1.0%, and Mo: 0.5%, the amount of hard martensite becomes excessive and the strength becomes higher than necessary. Therefore, when Cr, V and Mo are contained, Cr: 0.05% to 5.0%, V: 0.005% to 1.0%, and Mo: 0.005% to 0.00%. The range is 5% or less.
Ｔｉ：０．０１％以上０．１％以下、Ｎｂ：０．０１％以上０．１％以下のうちから選ばれる１種または２種
ＴｉおよびＮｂは鋼の析出強化に有用で、その効果は、それぞれの含有量が０.０１％以上で得られる。一方、それぞれの含有量が０．１％を超えると加工性および形状凍結性が低下する。従って、ＴｉおよびＮｂを含有させる場合は、Ｔｉ：０．０１％以上０．１％以下およびＮｂ：０．０１％以上０．１％以下の範囲とする。
One or two selected from Ti: 0.01% or more and 0.1% or less, Nb: 0.01% or more and 0.1% or less Ti and Nb are useful for the precipitation strengthening of steel. , Each content is obtained at 0.01% or more. On the other hand, when each content exceeds 0.1%, the workability and the shape freezing property are lowered. Therefore, when Ti and Nb are contained, the range is Ti: 0.01% to 0.1% and Nb: 0.01% to 0.1%.
Ｂ：０．０００３％以上０．００５０％以下
Ｂはオーステナイト粒界からフェライトが生成・成長することを抑制するのに有用な元素である。その効果は０．０００３％以上の含有で得られる。一方、含有量が０．００５０％を超えると加工性が低下する。従って、Ｂを含有させる場合は、Ｂ：０．０００３％以上０．００５０％以下の範囲とする。
B: 0.0003% or more and 0.0050% or less B is an element useful for suppressing the formation and growth of ferrite from the austenite grain boundary. The effect is obtained when the content is 0.0003% or more. On the other hand, if the content exceeds 0.0050%, the workability decreases. Therefore, when it contains B, it is set as B: 0.0003% or more and 0.0050% or less of range.
Ｎｉ：０．０５％以上２.０％以下およびＣｕ：０．０５％以上２.０％以下のうちから選ばれる１種または２種
ＮｉおよびＣｕは鋼の強化に有効な元素である。また、鋼板に溶融亜鉛めっきまたは合金化溶融亜鉛めっきを施す場合には、鋼板表層部の内部酸化を促進してめっき密着性を向上させる。これらの効果は、それぞれの含有量が０．０５％以上で得られる。一方、それぞれの含有量が２.０％を超えると、鋼板の加工性を低下させる．従って、ＮｉおよびＣｕを含有させる場合には、Ｎｉ：０．０５％以上２.０％以下およびＣｕ：０．０５％以上２.０％以下の範囲とする。
One or two selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less. Ni and Cu are effective elements for strengthening steel. Moreover, when performing hot dip galvanization or alloying hot dip galvanization to a steel plate, the internal oxidation of a steel plate surface layer part is accelerated | stimulated and plating adhesiveness is improved. These effects are obtained when the respective contents are 0.05% or more. On the other hand, if each content exceeds 2.0%, the workability of the steel sheet is lowered. Therefore, when Ni and Cu are contained, the range is Ni: 0.05% to 2.0% and Cu: 0.05% to 2.0%.
Ｃａ：０．００１％以上０．００５％以下およびＲＥＭ：０．００１％以上０．００５％以下のうちから選ばれる１種または２種
ＣａおよびＲＥＭは、硫化物の形状を球状化し、伸びフランジ性への硫化物の悪影響を改善するために有用である。その効果は、それぞれの含有量が０.００１％以上で得られる。一方、それぞれの含有量が０．００５％を超えると、介在物等の増加を招き、表面欠陥および内部欠陥などを引き起こす。従って、ＣａおよびＲＥＭを含有させる場合には、Ｃａ：０．００１％以上０．００５％以下およびＲＥＭ：０．００１％以上０．００５％以下の範囲とする。
One or two types selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less Ca and REM spheroidize the shape of the sulfide, and stretch flange Useful to improve the negative effects of sulfides on sex. The effect is obtained when each content is 0.001% or more. On the other hand, if the respective contents exceed 0.005%, inclusions and the like increase, causing surface defects and internal defects. Therefore, when Ca and REM are contained, the range is Ca: 0.001% to 0.005% and REM: 0.001% to 0.005%.
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。 In the steel plate 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、本発明の高強度鋼板の製造方法について説明する。
上記の好適成分組成に調整した鋼片を製造後、熱間圧延し、ついで冷間圧延を施して冷延鋼板とする。本発明において、これらの処理に特に制限はなく、常法に従って行えば良い。
好適な製造条件は次のとおりである。鋼片を、１０００℃以上１３００℃以下の温度域に加熱した後、８７０℃以上９５０℃以下の温度域で熱間圧延を終了し、得られた熱延鋼板を３５０℃以上７２０℃以下の温度域で巻き取る。ついで、熱延鋼板を酸洗後、４０％以上９０％以下の範囲の圧下率で冷間圧延を行い冷延鋼板とする。
なお、本発明では、鋼板を通常の製鋼、鋳造、熱間圧延、酸洗および冷間圧延の各工程を経て製造する場合を想定しているが、例えば、薄スラブ鋳造やストリップ鋳造などにより熱間圧延工程の一部または全部を省略して製造しても良い。
Next, the manufacturing method of the high strength steel plate of this invention is demonstrated.
A steel slab adjusted to the above preferred component composition is manufactured, then hot-rolled, and then cold-rolled to obtain a cold-rolled steel sheet. In the present invention, these treatments are not particularly limited, and may be performed according to ordinary methods.
The preferred production conditions are as follows. After heating the steel slab to a temperature range of 1000 ° C. or higher and 1300 ° C. or lower, hot rolling is finished in a temperature range of 870 ° C. or higher and 950 ° C. or lower, and the obtained hot rolled steel sheet is heated to a temperature of 350 ° C. or higher and 720 ° C. or lower. Take up in the area. Next, the hot-rolled steel sheet is pickled and then cold-rolled at a rolling reduction in the range of 40% to 90% to obtain a cold-rolled steel sheet.
In the present invention, it is assumed that the steel sheet is manufactured through normal steelmaking, casting, hot rolling, pickling and cold rolling processes. However, for example, the steel plate is heated by thin slab casting or strip casting. You may manufacture by omitting a part or all of a hot rolling process.
得られた冷延鋼板に、図１に示す熱処理を施す。以下、図１を参照しながら説明する。
オーステナイト単相域で１５秒以上６００秒以下の焼鈍をする。本発明の鋼板は、上部ベイナイトやマルテンサイトといった未変態オーステナイトから変態させて得る低温変態相を主相とするものであり、ポリゴナルフェライトは極力少ない方が好ましく、このためオーステナイト単相域での焼鈍が必要である。焼鈍温度に関しては、オーステナイト単相域であれば特に制限はないが、焼鈍温度が１０００℃を超えるとオーステナイト粒の成長が著しく、後の冷却によって生じる構成相の粗大化を引き起こし、靭性などを劣化させる。一方、焼鈍温度がＡ３点（オーステナイト変態点）未満の場合には、焼鈍段階で既にポリゴナルフェライトが生成しており、冷却中のポリゴナルフェライトの成長を抑制するためには５００℃以上の温度域を極めて急速に冷却する必要が生じる。従って、焼鈍温度は、Ａ３点（オーステナイト変態点）℃以上とする必要があり、１０００℃以下とすることが好ましい。
また、焼鈍時間が１５秒未満の場合には、オーステナイトへの逆変態が十分に進まない場合や、鋼板中の炭化物が十分に溶解しない場合がある。一方、焼鈍時間が６００秒を超えると、多大なエネルギー消費に伴うコスト増を招く。従って、焼鈍時間は１５秒以上６００秒以下の範囲とする。好ましくは、６０秒以上５００秒以下の範囲である。ここで、Ａ３点は、
Ａ３点(℃)＝９１０−２０３×[Ｃ％]1/2＋４４．７×[Ｓｉ％]−３０×[Ｍｎ％]
＋７００×[Ｐ％]＋１３０×[Ａｌ％]−１５．２×[Ｎｉ％]
−１１×[Ｃｒ％]−２０×[Ｃｕ％]＋３１．５×[Ｍｏ％]
＋１０４×[Ｖ％]＋４００×[Ｔｉ％]
によって算出することができる。なお、[Ｘ％]は、鋼板の成分元素Ｘの質量％とする。
The obtained cold-rolled steel sheet is subjected to the heat treatment shown in FIG. Hereinafter, a description will be given with reference to FIG.
Annealing is performed for 15 seconds to 600 seconds in an austenite single phase region. The steel sheet of the present invention is mainly composed of a low-temperature transformation phase obtained by transformation from untransformed austenite such as upper bainite and martensite. Polygonal ferrite is preferably as little as possible, and therefore, in the austenite single-phase region. Annealing is required. The annealing temperature is not particularly limited as long as it is in the austenite single phase region, but if the annealing temperature exceeds 1000 ° C., the growth of austenite grains is remarkable, which causes coarsening of the constituent phases caused by subsequent cooling, and deteriorates toughness and the like. Let On the other hand, the annealing temperature is of less than A 3 point (austenitic transformation point) is already generated by the polygonal ferrite in the annealing step, in order to suppress the growth of polygonal ferrite during cooling than 500 ° C. It becomes necessary to cool the temperature range very rapidly. Therefore, the annealing temperature must be the A 3 point (austenitic transformation point) ° C. or higher, preferably set to 1000 ° C. or less.
Moreover, when annealing time is less than 15 second, the reverse transformation to austenite may not fully advance, or the carbide | carbonized_material in a steel plate may not fully melt | dissolve. On the other hand, if the annealing time exceeds 600 seconds, the cost increases due to the great energy consumption. Accordingly, the annealing time is in the range of 15 seconds to 600 seconds. Preferably, it is the range of 60 seconds or more and 500 seconds or less. Here, A 3 points are
A 3 points (° C.) = 910−203 × [C%] 1/2 + 44.7 × [Si%] − 30 × [Mn%]
+ 700 × [P%] + 130 × [Al%] − 15.2 × [Ni%]
−11 × [Cr%] − 20 × [Cu%] + 31.5 × [Mo%]
+ 104 × [V%] + 400 × [Ti%]
Can be calculated. [X%] is mass% of the component element X of the steel sheet.
焼鈍後の冷延鋼板は、３５０℃以上４９０℃以下の保持温度域で定める冷却停止温度：Ｔ℃まで冷却されるが、少なくとも５５０℃までは、平均冷却速度を５℃/ｓ以上に制御して冷却される。平均冷却速度が５℃/s未満の場合、ポリゴナルフェライトの過剰な生成・成長やパーライト等の析出が生じ、所望の組織が得られない。従って、平均冷却速度は５℃/ｓ以上とする。好ましくは１０℃/ｓ以上である。平均冷却速度の上限は、冷却停止温度にバラツキが生じない限り特に限定されないが、一般的な設備では、平均冷却速度が１００℃/ｓを超えると、鋼板の長手方向および板幅方向での組織のバラツキが著しく大きくなるため、１００℃/ｓ以下が好ましい。従って、平均冷却速度は、１０℃/ｓ以上１００℃/ｓ以下の範囲が好ましい。 The cold-rolled steel sheet after annealing is cooled to a cooling stop temperature: T ° C. determined in a holding temperature range of 350 ° C. or higher and 490 ° C. or lower, but at least up to 550 ° C., the average cooling rate is controlled to 5 ° C./s or higher. And cooled. When the average cooling rate is less than 5 ° C./s, excessive formation and growth of polygonal ferrite and precipitation of pearlite occur, and a desired structure cannot be obtained. Therefore, the average cooling rate is 5 ° C./s or more. Preferably it is 10 degrees C / s or more. The upper limit of the average cooling rate is not particularly limited as long as the cooling stop temperature does not vary. In general equipment, when the average cooling rate exceeds 100 ° C./s, the structure in the longitudinal direction and the sheet width direction of the steel sheet. The dispersion is significantly increased, so that it is preferably 100 ° C./s or less. Therefore, the average cooling rate is preferably in the range of 10 ° C./s to 100 ° C./s.
５５０℃まで冷却された鋼板は、冷却停止温度：Ｔ℃まで引き続き冷却される。Ｔ℃以上５５０℃以下の温度域において鋼板が冷却される速度は、該保持温度域での保持時間を１５秒以上１０００秒以下とする以外は特に制限されないが、鋼板が過度に遅い速度で冷却された場合には、未変態オーステナイトから炭化物が生成することにより、所望の組織が得られなくなる可能性が高い。従って、Ｔ℃以上５５０℃以下の温度域において、鋼板は、平均で０.５℃/s以上の速度で冷却されることが好ましい。 The steel sheet cooled to 550 ° C. is continuously cooled to a cooling stop temperature: T ° C. The speed at which the steel sheet is cooled in the temperature range of T ° C. or more and 550 ° C. or less is not particularly limited except that the holding time in the holding temperature range is 15 seconds or more and 1000 seconds or less, but the steel sheet is cooled at an excessively slow speed. In such a case, there is a high possibility that a desired structure cannot be obtained due to the formation of carbide from untransformed austenite. Therefore, in the temperature range of T ° C. or higher and 550 ° C. or lower, the steel plate is preferably cooled at an average rate of 0.5 ° C./s or higher.
冷却停止温度：Ｔ℃まで冷却された鋼板は、４９０℃からＴ℃までの冷却に要する時間を含め３５０〜４９０℃の保持温度域で１５秒以上１０００秒以下の時間保持される。すなわち該保持時間は、焼鈍後の鋼板が３５０℃以上４９０℃以下の範囲にある全時間である。保持温度域の上限が４９０℃を超えると、未変態オーステナイトから炭化物が析出して、本発明で規定する組織が得られない場合がある。一方、保持温度域の下限が３５０℃未満の場合、上部ベイナイトではなく、下部ベイナイトが生成し、オーステナイト中へのＣ濃化量が少なくなることが問題となる。従って、保持温度域の範囲は、３５０℃以上４９０℃以下の範囲とする。好ましくは、３７０℃以上４６０℃以下の範囲である。
また、保持温度域での保持時間が１５秒未満の場合、ベイナイト変態が十分に進行せず、所望の組織が得られない。一方、保持温度域での保持時間が１０００秒を超えると、鋼板の最終組織として残留オーステナイトとなる未変態オーステナイトから炭化物が析出してＣが濃化した安定な残留オーステナイトが得られず、その結果、所望の強度と延性またはその両方が得られない。従って、保持時間は１５秒以上１０００秒以下とする。好ましくは、３０秒以上６００秒以下の範囲である。さらに好ましくは、６０秒以上４５０秒以下である。
Cooling stop temperature: The steel sheet cooled to T ° C. is held for 15 seconds to 1000 seconds in a holding temperature range of 350 to 490 ° C. including the time required for cooling from 490 ° C. to T ° C. That is, the holding time is the total time that the steel sheet after annealing is in the range of 350 ° C. or higher and 490 ° C. or lower. If the upper limit of the holding temperature range exceeds 490 ° C., carbide may precipitate from untransformed austenite, and the structure defined in the present invention may not be obtained. On the other hand, when the lower limit of the holding temperature range is less than 350 ° C., lower bainite is generated instead of upper bainite, and the amount of C enrichment in austenite becomes a problem. Accordingly, the range of the holding temperature range is 350 ° C. or more and 490 ° C. or less. Preferably, it is the range of 370 degreeC or more and 460 degreeC or less.
In addition, when the holding time in the holding temperature range is less than 15 seconds, the bainite transformation does not proceed sufficiently and a desired structure cannot be obtained. On the other hand, if the holding time in the holding temperature range exceeds 1000 seconds, stable residual austenite in which carbide is precipitated and C is concentrated from the untransformed austenite which becomes the retained austenite as the final structure of the steel sheet cannot be obtained. The desired strength and ductility or both cannot be obtained. Therefore, the holding time is 15 seconds or more and 1000 seconds or less. Preferably, it is the range of 30 seconds or more and 600 seconds or less. More preferably, it is 60 seconds or more and 450 seconds or less.
なお、本発明における一連の熱処理では、上述した所定の温度範囲内であれば、保持温度は一定である必要はなく、所定の温度範囲内で変動しても本発明の趣旨を損なわない。冷却速度についても同様である。また、熱履歴さえ満足すれば、鋼板はいかなる設備で熱処理を施されても構わない。さらに、熱処理後に、形状矯正のために鋼板の表面に調質圧延を施すことや電気めっき等の表面処理を施すことも本発明の範囲に含まれる。 In the series of heat treatments according to the present invention, the holding temperature does not have to be constant as long as it is within the predetermined temperature range described above, and the gist of the present invention is not impaired even if it fluctuates within the predetermined temperature range. The same applies to the cooling rate. Further, as long as the thermal history is satisfied, the steel sheet may be heat-treated with any equipment. Further, it is also included in the scope of the present invention that after the heat treatment, the surface of the steel sheet is subjected to temper rolling or surface treatment such as electroplating for shape correction.
本発明の高強度鋼板の製造方法には、さらに、溶融亜鉛めっき処理、あるいは溶融亜鉛めっき処理にさらに合金化処理を加えた合金化溶融亜鉛めっき処理を加えることができる。溶融亜鉛めっき処理、あるいは合金化溶融亜鉛めっき処理は、上記した保持温度域までの冷却中もしくは保持温度域で行っても良い。この場合、保持温度域での保持時間は、溶融亜鉛めっき処理あるいは合金化亜鉛めっき処理の保持温度域における処理時間も含めて１５秒以上１０００秒以下とする。なお、該溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理は、連続溶融亜鉛めっきラインにて行うことが好ましい。 The method for producing a high-strength steel sheet according to the present invention can further include hot dip galvanizing treatment or galvannealing treatment in which alloying treatment is further added to hot dip galvanizing treatment. The hot dip galvanizing process or the alloyed hot dip galvanizing process may be performed during cooling to the holding temperature range or in the holding temperature range. In this case, the holding time in the holding temperature range is 15 seconds or more and 1000 seconds or less including the processing time in the holding temperature range of the hot dip galvanizing treatment or alloying galvanizing treatment. The hot dip galvanizing treatment or alloying hot dip galvanizing treatment is preferably performed in a continuous hot dip galvanizing line.
また、本発明の高強度鋼板の製造方法では、上記した本発明の製造方法に従い、熱処理まで完了した高強度鋼板を製造した後、改めて溶融亜鉛めっき処理、あるいはさらに合金化処理を施すことを加えることができる。 Moreover, in the manufacturing method of the high-strength steel sheet according to the present invention, after the high-strength steel sheet completed up to the heat treatment is manufactured according to the above-described manufacturing method of the present invention, a hot dip galvanizing process or a further alloying process is added. be able to.
鋼板に溶融亜鉛めっき処理または合金化溶融亜鉛めっき処理を行う方法は、次のとおりである。
鋼板をめっき浴中に浸入させ、ガスワイピングなどで付着量を調整する。めっき浴中の溶解Ａｌ量は、溶融亜鉛めっき処理の場合は０．１２％以上０．２２％以下の範囲、合金化溶融亜鉛めっき処理の場合は０．０８％以上０．１８％以下の範囲とすることが好ましい。
処理温度は、溶融亜鉛めっき処理の場合、めっき浴の温度は通常の４５０℃以上５００℃以下の範囲であればよく、さらに合金化処理を施す場合、合金化時の温度は５５０℃以下とすることが好ましい。合金化温度が５５０℃を超える場合、未変態オーステナイトから炭化物が析出したり、場合によってはパーライトが生成するため、強度や加工性またはその両方が得られず、また、めっき層のパウダリング性も劣化する。一方、合金化時の温度が４５０℃未満では合金化が進行しない場合があるため、４５０℃以上とすることが好ましい。
めっき付着量は片面当たり２０ｇ/ｍ２以上１５０ｇ/ｍ２以下の範囲とすることが好ましい。めっき付着量が２０ｇ/ｍ２未満では耐食性が不足し、一方、１５０ｇ/ｍ２を超えても耐食効果は飽和し、コストアップを招くだけである。
めっき層の合金化度（Ｆｅ質量％(Ｆｅ含有量)）は７質量％以上１５質量％以下の範囲が好ましい。めっき層の合金化度が７質量％未満では、合金化ムラが生じ外観品質が劣化したり、めっき層中にいわゆるζ相が生成され鋼板の摺動性が劣化したりする。一方、めっき層の合金化度が１５質量％を超えると、硬質で脆いΓ相が多量に形成され、めっき密着性が劣化する。
The method of performing hot dip galvanizing treatment or alloying hot dip galvanizing treatment on a steel sheet is as follows.
The steel sheet is infiltrated into the plating bath and the amount of adhesion is adjusted by gas wiping. The amount of dissolved Al in the plating bath ranges from 0.12% to 0.22% in the case of hot dip galvanizing, and ranges from 0.08% to 0.18% in the case of galvannealed alloying. It is preferable that
In the case of hot dip galvanizing, the temperature of the plating bath may be in the range of 450 ° C. or higher and 500 ° C. or lower. When further alloying is performed, the temperature during alloying is 550 ° C. or lower. It is preferable. When the alloying temperature exceeds 550 ° C., carbide precipitates from untransformed austenite or pearlite is generated in some cases, so that strength and workability or both cannot be obtained, and the powdering property of the plating layer is also low. to degrade. On the other hand, if the temperature during alloying is less than 450 ° C., alloying may not proceed.
Coating weight is preferably in a per side 20 g / m 2 or more 150 g / m 2 or less. If the plating adhesion amount is less than 20 g / m 2 , the corrosion resistance is insufficient. On the other hand, if it exceeds 150 g / m 2 , the corrosion resistance effect is saturated and only the cost is increased.
The alloying degree (Fe mass% (Fe content)) of the plating layer is preferably in the range of 7 mass% to 15 mass%. If the degree of alloying of the plating layer is less than 7% by mass, unevenness in alloying occurs and the appearance quality deteriorates, or the so-called ζ phase is generated in the plating layer and the slidability of the steel sheet deteriorates. On the other hand, when the degree of alloying of the plating layer exceeds 15% by mass, a large amount of hard and brittle Γ phase is formed, and the plating adhesion deteriorates.
以下、本発明を実施例によってさらに詳細に説明するが、下記実施例は本発明を限定するものではない。また、本発明の要旨構成の範囲内で構成を変更することは、本発明の範囲に含まれるものとする。 EXAMPLES Hereinafter, although an Example demonstrates this invention further in detail, the following Example does not limit this invention. In addition, changing the configuration within the scope of the gist configuration of the present invention is included in the scope of the present invention.
表１に示す成分組成の鋼を溶製して得た鋳片を、１２００℃に加熱し、８７０℃で仕上げ熱間圧延した熱延鋼板を６５０℃で巻き取り、ついで熱延鋼板を酸洗後、６５％の圧延率（圧下率）で冷間圧延し、板厚：１.２ｍｍの冷延鋼板とした。得られた冷延鋼板を、表２に示す条件で熱処理を施した。
また、一部の冷延鋼板については、溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理を施した。ここで、溶融亜鉛めっき処理は、めっき浴温度：４６３℃、目付け量（片面あたり）：５０ｇ/ｍ２となるように両面めっきを施した。また、合金化溶融亜鉛めっき処理は、同じくめっき浴温度：４６３℃、目付け量（片面あたり）：５０ｇ/ｍ２として合金化度（Ｆｅ質量％(Ｆｅ含有量)）が９質量％となるように合金化温度：５５０℃以下で合金化条件を調整して両面めっきを施した。なお、溶融亜鉛めっき処理および合金化溶融亜鉛めっき処理は、表２中に示すＴ℃まで一旦冷却した後に行った。
The slab obtained by melting the steel having the composition shown in Table 1 is heated to 1200 ° C, the hot-rolled steel sheet finished by hot rolling at 870 ° C is wound up at 650 ° C, and then the hot-rolled steel sheet is pickled. Thereafter, it was cold-rolled at a rolling rate (rolling rate) of 65% to obtain a cold-rolled steel plate having a thickness of 1.2 mm. The obtained cold-rolled steel sheet was heat-treated under the conditions shown in Table 2.
Further, some cold-rolled steel sheets were subjected to hot dip galvanizing treatment or alloying hot dip galvanizing treatment. Here, in the hot dip galvanizing treatment, double-sided plating was performed so that the plating bath temperature was 463 ° C. and the basis weight (per one side) was 50 g / m 2 . Similarly, the alloying hot dip galvanizing treatment is performed so that the plating bath temperature is 463 ° C., the basis weight (per one side) is 50 g / m 2 , and the degree of alloying (Fe mass% (Fe content)) is 9 mass%. The alloying temperature was adjusted to 550 ° C. or lower and the alloying conditions were adjusted to perform double-sided plating. In addition, the hot dip galvanizing treatment and the alloying hot dip galvanizing treatment were performed after cooling to T ° C shown in Table 2 once.
得られた鋼板に、めっき処理を施さない場合には熱処理後に、溶融亜鉛めっき処理あるいは合金化溶融亜鉛めっき処理を施す場合にはこれらの処理後に、圧延率（伸び率）：０．３％の調質圧延を施した。 When the obtained steel sheet is not subjected to plating treatment, after heat treatment, when subjected to hot dip galvanizing treatment or alloyed hot dip galvanizing treatment, after these treatments, the rolling rate (elongation rate): 0.3% Temper rolling was performed.
かくして得られた鋼板の諸特性を以下の方法で評価した。
各鋼板から試料を切り出し研磨して、圧延方向に平行な面を走査型電子顕微鏡（ＳＥＭ）を用いて３０００倍で１０視野組織観察して、各相の面積率を測定し、各結晶粒の相構造を同定した。
Various properties of the steel sheet thus obtained were evaluated by the following methods.
Samples were cut from each steel plate and polished, and the surface parallel to the rolling direction was observed with a scanning electron microscope (SEM) at 10Ox field of view at 10 000 times to measure the area ratio of each phase. The phase structure was identified.
残留オーステナイト量は、鋼板を板厚方向に板厚の１/４まで研削・研磨し、Ｘ線回折強度測定により求めた。入射Ｘ線には、Ｃｏ−Ｋαを用い、フェライトの(２００)、(２１１)、(２２０)各面の回折強度に対するオーステナイトの(２００)、(２２０)、（３１１）各面の強度比から残留オーステナイト量を計算した。 The amount of retained austenite was determined by measuring the X-ray diffraction intensity after grinding and polishing the steel plate to ¼ of the plate thickness in the plate thickness direction. For incident X-rays, Co—Kα is used, and from the intensity ratio of each surface of austenite (200), (220), (311) to the diffraction intensity of each surface of ferrite (200), (211), (220). The amount of retained austenite was calculated.
残留オーステナイト中の平均Ｃ量は、Ｘ線回折強度測定でのオーステナイトの(２００)、(２２０)、(３１１)各面の強度ピークから格子定数を求め、次の計算式から残留オーステナイト中の平均Ｃ量（質量％）を求めた。
ａ０＝０．３５８０＋０．００３３×[Ｃ％]＋０．０００９５×[Ｍｎ％]
＋０．００５６×[Ａｌ％]＋０．０２２×[Ｎ％]
ただし、ａ０：格子定数（ｎｍ）、[Ｘ％]：元素Ｘの質量％。なお、Ｃ以外の元素の質量％は、鋼板全体に対する質量％とした。
The average amount of C in the retained austenite is obtained by calculating the lattice constant from the intensity peaks of the (200), (220) and (311) surfaces of austenite in the X-ray diffraction intensity measurement. C amount (mass%) was calculated | required.
a 0 = 0.3580 + 0.0033 × [C%] + 0.00095 × [Mn%]
+ 0.0056 × [Al%] + 0.022 × [N%]
However, a 0: the lattice constant (nm), [X%] : % by weight of the element X. In addition, mass% of elements other than C was mass% with respect to the whole steel plate.
引張試験は、鋼板の圧延方向に対して垂直な方向から採取したＪＩＳ５号試験片を用いて、ＪＩＳ Ｚ２２４１に準拠して行った。ＴＳ（引張強さ）、Ｔ.Ｅｌ（全伸び）を測定し、強度と全伸びの積（ＴＳ×Ｔ．Ｅｌ）を算出して、加工性（延性）を評価した。なお、本発明では、ＴＳ×Ｔ.Ｅｌ≧２００００（ＭＰａ・％）の場合を良好とした。 The tensile test was performed according to JIS Z2241, using a JIS No. 5 test piece taken from a direction perpendicular to the rolling direction of the steel sheet. TS (tensile strength) and T.El (total elongation) were measured, and the product of strength and total elongation (TS × T.El) was calculated to evaluate workability (ductility). In the present invention, the case of TS × T.E1 ≧ 20000 (MPa ·%) was considered good.
以上の評価結果を表３に示す。 The above evaluation results are shown in Table 3.
同表から明らかなように、本発明の鋼板はいずれも、引張強さが９８０ＭＰａ以上、かつＴＳ×Ｔ.Ｅｌの値が２００００ＭＰａ・％以上を満足することから、高強度と優れた加工性を兼ね備えていることが確認できた。 As is apparent from the table, all the steel plates of the present invention have a tensile strength of 980 MPa or more and a value of TS × T.El of 20000 MPa ·% or more, and thus have high strength and excellent workability. It was confirmed that they had both.
これに対し、試料Ｎｏ．１は５５０℃までの平均冷却速度Ａ１が５℃/ｓ未満であることから、所望の鋼板組織が得られず、引張強さ（ＴＳ）：９８０ＭＰａ以上およびＴＳ×Ｔ.Ｅｌの値が２００００ＭＰａ・％以上の両方を満足せず、強度、加工性ともに劣るものであった。試料Ｎｏ．３は保持温度域での保持時間が１５秒未満であることから、試料Ｎｏ．４は焼鈍温度がＡ３点未満であることから、試料Ｎｏ．６は焼鈍時間が１５秒未満であることから、所望の鋼板組織が得られず、その結果、引張強さ（ＴＳ）は９８０ＭＰａ以上を満足するものの、ＴＳ×Ｔ.Ｅｌの値が２００００ＭＰａ・％未満で加工性に劣る。試料Ｎｏ．２９および３０は、成分組成が本発明の適正範囲外であることから、所望の鋼板組織が得られず、その結果、引張強さ（ＴＳ）は９８０ＭＰａ以上を満足するものの、ＴＳ×Ｔ.Ｅｌの値が２００００ＭＰａ・％未満で加工性に劣る。試料Ｎｏ．３１および３２は、成分組成が本発明の適正範囲外であることから、所望の鋼板組織が得られず、引張強さ（ＴＳ）：９８０ＭＰａ以上およびＴＳ×Ｔ.Ｅｌの値が２００００ＭＰａ・％以上の両方を満足せず、強度、加工性ともに劣るものであった。 In contrast, sample no. No. 1 has an average cooling rate A1 up to 550 ° C. of less than 5 ° C./s, so that a desired steel sheet structure cannot be obtained. Tensile strength (TS): 980 MPa or more and TS × T.El is 20000 MPa · % Was not satisfied, and both strength and workability were inferior. Sample No. 3 has a holding time in the holding temperature range of less than 15 seconds. 4 since the annealing temperature is less than 3 points A, Sample No. No. 6 has an annealing time of less than 15 seconds, so a desired steel sheet structure cannot be obtained. As a result, although the tensile strength (TS) satisfies 980 MPa or more, the value of TS × T.El is 20000 MPa ·%. Less than is inferior in workability. Sample No. In Nos. 29 and 30, since the component composition is outside the proper range of the present invention, a desired steel sheet structure cannot be obtained. As a result, although the tensile strength (TS) satisfies 980 MPa or more, TS × T.El Is less than 20000 MPa ·%, the workability is poor. Sample No. 31 and 32, since the component composition is outside the proper range of the present invention, the desired steel sheet structure cannot be obtained, and the tensile strength (TS): 980 MPa or more and the value of TS × T.El is 20000 MPa ·% or more. Both were not satisfied, and both strength and workability were inferior.
Claims (9)
Ｃ：０.１６％以上０.７２％以下、
Ｓｉ：３.０％以下、
Ｍｎ：０.５％以上３.０％以下、
Ｐ：０．１％以下、
Ｓ：０．０７％以下、
Ａｌ：３.０％以下および
Ｎ：０．０１０％以下
を含有し、かつＳｉ＋Ａｌが０.７％以上を満足し、残部はＦｅおよび不可避不純物の組成になり、
鋼板組織として、マルテンサイトの鋼板組織全体に対する面積率が１０％以上９０％以下、残留オーステナイト量が５％以上５０％以下、上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率が５％以上であり、前記マルテンサイトの鋼板組織全体に対する面積率、前記残留オーステナイト量および前記上部ベイナイト中のベイニティックフェライトの鋼板組織全体に対する面積率の合計が６０％以上、ポリゴナルフェライトの鋼板組織全体に対する面積率が１０％以下（０％を含む）を満足し、かつ前記残留オーステナイト中の平均Ｃ量が０.７０％以上２.００％以下であり、引張強さが９８０ＭＰａ以上であることを特徴とする高強度鋼板。 C: 0.16% to 0.72% by mass%,
Si: 3.0% or less,
Mn: 0.5% to 3.0%,
P: 0.1% or less,
S: 0.07% or less,
Al: 3.0% or less and N: 0.010% or less, and Si + Al satisfies 0.7% or more, the balance is the composition of Fe and inevitable impurities,
As the steel sheet structure, the area ratio of martensite to the entire steel sheet structure is 10% to 90%, the amount of retained austenite is 5% to 50%, and the area ratio of bainitic ferrite in the upper bainite to the entire steel sheet structure is 5%. The total area ratio of the martensite to the entire steel sheet structure, the amount of retained austenite, and the area ratio of the bainitic ferrite in the upper bainite to the entire steel sheet structure is 60% or more, and the entire steel structure of polygonal ferrite That the area ratio with respect to the above is 10% or less (including 0%), the average C content in the retained austenite is 0.70% or more and 2.00% or less, and the tensile strength is 980 MPa or more. High-strength steel sheet characterized.
Ｃｒ：０.０５％以上５.０％以下、
Ｖ：０．００５％以上１．０％以下および
Ｍｏ：０．００５％以上０．５％以下
のうちから選ばれる１種または２種以上の元素を含有することを特徴とする請求項１に記載の高強度鋼板。 The steel sheet is further in mass%,
Cr: 0.05% to 5.0%,
The element according to claim 1, comprising one or more elements selected from V: 0.005% to 1.0% and Mo: 0.005% to 0.5%. High strength steel sheet as described.
Ｔｉ：０．０１％以上０．１％以下および
Ｎｂ：０．０１％以上０．１％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする請求項１または２に記載の高強度鋼板。 The steel sheet is further in mass%,
3. One or two elements selected from Ti: 0.01% or more and 0.1% or less and Nb: 0.01% or more and 0.1% or less are contained. The high-strength steel sheet described in 1.
Ｂ：０．０００３％以上０．００５０％以下
を含有することを特徴とする請求項１乃至３のいずれか１項に記載の高強度鋼板。 The steel sheet is further in mass%,
The high-strength steel plate according to any one of claims 1 to 3, wherein B: 0.0003% or more and 0.0050% or less is contained.
Ｎｉ：０．０５％以上２.０％以下および
Ｃｕ：０．０５％以上２.０％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする請求項１乃至４のいずれか１項に記載の高強度鋼板。 The steel sheet is further in mass%,
5. One or two elements selected from Ni: 0.05% or more and 2.0% or less and Cu: 0.05% or more and 2.0% or less are contained. The high-strength steel sheet according to any one of the above.
Ｃａ：０．００１％以上０．００５％以下および
ＲＥＭ：０．００１％以上０．００５％以下
のうちから選ばれる１種または２種の元素を含有することを特徴とする請求項１乃至５のいずれか１項に記載の高強度鋼板。 The steel sheet is further in mass%,
6. One or two elements selected from Ca: 0.001% or more and 0.005% or less and REM: 0.001% or more and 0.005% or less are contained. The high-strength steel sheet according to any one of the above.
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