JP2009185355A - High strength cold-rolled steel sheet having excellent workability and collision resistance and its production method - Google Patents
High strength cold-rolled steel sheet having excellent workability and collision resistance and its production method Download PDFInfo
- Publication number
- JP2009185355A JP2009185355A JP2008027856A JP2008027856A JP2009185355A JP 2009185355 A JP2009185355 A JP 2009185355A JP 2008027856 A JP2008027856 A JP 2008027856A JP 2008027856 A JP2008027856 A JP 2008027856A JP 2009185355 A JP2009185355 A JP 2009185355A
- Authority
- JP
- Japan
- Prior art keywords
- steel sheet
- ferrite
- rolled steel
- cold
- workability
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Granted
Links
Images
Landscapes
- Metal Rolling (AREA)
- Coating With Molten Metal (AREA)
- Heat Treatment Of Sheet Steel (AREA)
Abstract
Description
本発明は、自動車用鋼板等の用途に好適な高強度冷延鋼板及びその製造方法に関するものである。 The present invention relates to a high-strength cold-rolled steel sheet suitable for uses such as automotive steel sheets and a method for producing the same.
炭酸ガスの排出量の抑制するため、自動車の燃費の向上を目的とする自動車車体の軽量化が進められている。そのため、自動車の部材には、板厚の低減が可能な高強度鋼板の適用が増えつつある。また、搭乗者の安全性確保のためにも、高強度鋼板が自動車車体に多く使用されるようになってきている。 In order to reduce the amount of carbon dioxide emissions, the weight reduction of automobile bodies for the purpose of improving the fuel efficiency of automobiles is being promoted. Therefore, the application of high-strength steel sheets capable of reducing the plate thickness is increasing for automobile members. Further, in order to ensure the safety of passengers, high-strength steel plates are increasingly used in automobile bodies.
一方、高強度鋼板を自動車車体に適用するためには優れた加工性も要求される。このような強度と加工性を両立させた鋼材として、フェライトとマルテンサイトを主体とする硬質第２相からなる複合組織を有する二相組織鋼（Ｄｕａｌ Ｐｈａｓｅ鋼、以下、ＤＰ鋼）が知られている。しかしながら、ＤＰ鋼は主相を軟質なフェライトとし、硬質組織であるベイナイト、マルテンサイト、焼き戻しマルテンサイトを用いて強化していることから、降伏比が低い。したがって、部材としての変形を抑えつつ、衝突エネルギーを吸収する部材へのＤＰ鋼の適用には課題があった。 On the other hand, in order to apply a high-strength steel sheet to an automobile body, excellent workability is also required. As a steel material having both such strength and workability, a dual-phase steel (Dual Phase steel, hereinafter referred to as DP steel) having a composite structure composed of a hard second phase mainly composed of ferrite and martensite is known. Yes. However, DP steel has a low yield ratio because DP steel is made of soft ferrite and strengthened by using hard structures such as bainite, martensite, and tempered martensite. Therefore, there has been a problem in applying DP steel to a member that absorbs collision energy while suppressing deformation as a member.
また、フェライトの再結晶を防止し、未再結晶フェライトと硬質第２相からなる組織を有する鋼板が提案されている（例えば、特許文献１）。しかし、フェライトを未再結晶フェライトにすると、強度及び降伏比が高くなるものの、伸びが低いため、成形性が不十分であった。このような問題に対して、結晶粒の微細化や、析出強化とフェライト中の固溶Ｃ量の低減によって、フェライトとパーライトからなる組織を有する鋼板の高強度化と伸びフランジ性の向上の両立を図った鋼板が提案されている（例えば、特許文献２〜４）。しかし、引張強度がいずれも５００ＭＰａ以下であり、５９０ＭＰａを超えるような高強度化は困難であった。
In addition, a steel sheet that prevents recrystallization of ferrite and has a structure composed of non-recrystallized ferrite and a hard second phase has been proposed (for example, Patent Document 1). However, when the ferrite is non-recrystallized ferrite, the strength and the yield ratio are increased, but the elongation is low, and the formability is insufficient. For such problems, both strengthening of steel sheets with a structure composed of ferrite and pearlite and improvement of stretch flangeability are achieved by refining crystal grains, strengthening precipitation, and reducing the amount of dissolved C in ferrite. Have been proposed (for example,
これに対して、本発明者らの一部は、未再結晶フェライトの活用に着目し、軟質のフェライトと硬質第２相の中間の硬度を有する未再結晶フェライトにより、伸びフランジ性を向上させた鋼板を提案した（例えば、特許文献５〜１０）。 On the other hand, some of the present inventors pay attention to the utilization of non-recrystallized ferrite and improve stretch flangeability by using non-recrystallized ferrite having a hardness intermediate between soft ferrite and hard second phase. Steel plates were proposed (for example, Patent Documents 5 to 10).
しかし、特許文献５で提案されている鋼板は、ＮｂやＴｉの添加による析出強化を活用するものであり、パーライトの活用によって降伏比を高めるものではない。また、特許文献６及び７で提案されている鋼板は、ＮｂやＴｉの添加量が少なく、再結晶抑制効果が小さいため、焼鈍時に急速加熱を行う必要があった。また、特許文献８及び９で提案されている鋼板は、ＮｂやＴｉの添加量が少なく、昇温速度も１０℃／ｓ以下であったため、未再結晶フェライトの活用による効果が十分では無かったと考えられる。また、特許文献１０で提案されている鋼板は、Ｂが必須添加元素であり、合金コストがかかってしまうという問題があった。 However, the steel sheet proposed in Patent Document 5 utilizes precipitation strengthening due to the addition of Nb and Ti, and does not increase the yield ratio by utilizing pearlite. In addition, the steel plates proposed in Patent Documents 6 and 7 have a small amount of Nb and Ti added and have a small effect on suppressing recrystallization, and thus have to be rapidly heated during annealing. In addition, the steel plates proposed in Patent Documents 8 and 9 have a small amount of Nb and Ti added, and the rate of temperature increase is 10 ° C./s or less, so that the effect of utilizing non-recrystallized ferrite is not sufficient. Conceivable. Further, the steel sheet proposed in Patent Document 10 has a problem that B is an essential additive element and the alloy cost is increased.
本発明の課題は、成形性及び穴拡げ性に優れ、更に、降伏比が高く、耐衝突特性にも優れた高強度冷延鋼板の提供、及びこのような鋼板を安定的に生産性を損なうことなく得る製造方法の提供である。 An object of the present invention is to provide a high-strength cold-rolled steel sheet that is excellent in formability and hole expansibility, has a high yield ratio, and has excellent impact resistance characteristics, and stably impairs the productivity of such a steel sheet. It is providing the manufacturing method obtained without it.
本発明者らは、未再結晶フェライトを積極的に活用し、更に、マルテンサイト、ベイナイトの生成を制限して、フェライトとパーライトからなる金属組織とすることにより、降伏比が高く、且つ成形性及び穴拡げ性にも優れた高強度冷延鋼板を得ることができることを見出した。 The present inventors actively utilize non-recrystallized ferrite, further limit the formation of martensite and bainite, and form a metal structure composed of ferrite and pearlite, so that the yield ratio is high and the formability is high. And it discovered that the high-strength cold-rolled steel plate excellent also in the hole expansibility was able to be obtained.
また、本発明者らは、マルテンサイト、ベイナイトの生成を抑制して、フェライトとパーライトからなる金属組織とし、更に、未再結晶フェライトを残留させるための製造条件について検討した。その結果、焼入れ性の指標である炭素当量Ｃｅｑ、Ｎｂ及びＴｉの添加量の最適化に加えて、焼鈍工程における再結晶温度からＡｃ1変態温度までの昇温速度、更にフェライトとオーステナイトが共存する領域であるα＋γ二相域、即ち、Ａｃ1変態温度以上に加熱した際に、オーステナイトへの変態が進み過ぎないように、鋼板の温度がＡｃ1変態温度以上である滞留時間及び焼鈍の最高到達温度、その後の冷却速度の最適化が重要であることを見出した。 In addition, the present inventors studied production conditions for suppressing the formation of martensite and bainite to form a metal structure composed of ferrite and pearlite, and for leaving unrecrystallized ferrite. As a result, in addition to optimizing the addition amount of carbon equivalents Ceq, Nb and Ti, which is an index of hardenability, the heating rate from the recrystallization temperature to the Ac 1 transformation temperature in the annealing process, and further ferrite and austenite coexist. Α + γ two-phase region that is the region, that is, when the steel sheet is heated to a temperature higher than the Ac 1 transformation temperature, so that the transformation to austenite does not proceed excessively, the residence time when the temperature of the steel sheet is higher than the Ac 1 transformation temperature and the highest reach of annealing. We have found that optimization of temperature and subsequent cooling rate is important.
本発明は、このような知見に基づいてなされたものであり、その要旨は以下の通りである。 The present invention has been made based on such findings, and the gist thereof is as follows.
（１） 質量％で、Ｃ：０．０５〜０．２５％、Ｍｎ：０．５０〜２．５０％、を含有し、Ｓｉ：１．００％以下、Ａｌ：０．２００％以下、Ｐ：０．１５０％以下、Ｓ：０．０１５０％以下、Ｎ：０．０１００％以下に制限し、更に、Ｎｂ、Ｔｉの一方又は双方を合計で０．０４〜０．０８％含有し、下記（式１）によって求められるＣｅｑが０．４５以下であり、残部が鉄及び不可避的不純物からなり、金属組織がフェライトとパーライトからなり、前記フェライトが再結晶フェライト、変態フェライトの一方又は双方と未再結晶フェライトからなり、前記未再結晶フェライトの面積率が２０〜５０％であり、前記再結晶フェライト、前記変態フェライトの一方又は双方の面積率が２０〜７９％であり、前記パーライトの面積率が１〜３０％であることを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板。
Ｃｅｑ＝Ｃ＋Ｓｉ／２４＋Ｍｎ／６＋Ｎｉ／４０＋Ｃｒ／５＋Ｍｏ／４ ・・・（式１）
ここで、Ｃ、Ｓｉ、Ｍｎ、Ｎｉ、Ｃｒ、Ｍｏは、各元素の含有量［質量％］である。
(1) By mass%, C: 0.05 to 0.25%, Mn: 0.50 to 2.50%, Si: 1.00% or less, Al: 0.200% or less, P : 0.150% or less, S: 0.0150% or less, N: 0.0100% or less, and further contains one or both of Nb and Ti in a total amount of 0.04 to 0.08%. The Ceq obtained by (Equation 1) is 0.45 or less, the balance is made of iron and inevitable impurities, the metal structure is made of ferrite and pearlite, and the ferrite is uncrystallized with one or both of recrystallized ferrite and transformed ferrite. It consists of recrystallized ferrite, the area ratio of the unrecrystallized ferrite is 20 to 50%, the area ratio of one or both of the recrystallized ferrite and the transformed ferrite is 20 to 79%, and the area ratio of the pearlite 1-30 % High-strength cold-rolled steel sheet with excellent workability and impact resistance.
Ceq = C + Si / 24 + Mn / 6 + Ni / 40 + Cr / 5 + Mo / 4 (Formula 1)
Here, C, Si, Mn, Ni, Cr, and Mo are content [mass%] of each element.
（２） 冷延鋼板の降伏比が、０．７０以上であることを特徴とする上記（１）に記載の加工性及び耐衝突特性に優れた高強度冷延鋼板。
（３） 質量％で、Ｍｏ：０．１〜１．０％、Ｂ：０．０００５〜０．０１００％、Ｃｒ：０．１０〜１．５０％、Ｎｉ：０．１０〜１．５０％のうち、１種又は２種以上を含有することを特徴とする上記（１）又は（２）の何れか１項に記載の加工性及び耐衝突特性に優れた高強度冷延鋼板。
（４） 上記（１）〜（３）の何れか１項に記載の冷延鋼板の表面に溶融Ｚｎめっきを設けたことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板。
（５） 上記（１）〜（３）の何れか１項に記載の冷延鋼板の表面に合金化溶融Ｚｎめっきを設けたことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板。
(2) The high-strength cold-rolled steel sheet having excellent workability and impact resistance as described in (1) above, wherein the yield ratio of the cold-rolled steel sheet is 0.70 or more.
(3) By mass%, Mo: 0.1 to 1.0%, B: 0.0005 to 0.0100%, Cr: 0.10 to 1.50%, Ni: 0.10 to 1.50% Among them, the high-strength cold-rolled steel sheet having excellent workability and impact resistance according to any one of the above (1) or (2), comprising one or more of them.
(4) A high-strength cold-rolled steel sheet excellent in workability and impact resistance characteristics, characterized in that hot-dip Zn plating is provided on the surface of the cold-rolled steel sheet according to any one of (1) to (3) above .
(5) High-strength cooling excellent in workability and impact resistance characteristics, characterized by providing alloyed hot-dip Zn plating on the surface of the cold-rolled steel sheet according to any one of (1) to (3) above Rolled steel sheet.
（６） 上記（１）〜（３）の何れか１項に記載の化学成分を有する鋼片を熱間圧延し、酸洗後、冷間圧延を施した後、鋼板を、（Ａｃ1［℃］−１００℃）からＡｃ1［℃］までの昇温速度を５℃／ｓ以上としてＡｃ1［℃］〜｛Ａｃ1［℃］＋２／３×（Ａｃ3［℃］−Ａｃ1［℃］）｝の温度範囲内に昇温し、前記鋼板の温度が該温度範囲内である滞留時間を１０〜３００ｓとして焼鈍し、平均冷却速度を４０℃／ｓ以下として冷却することを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
ここで、Ａｃ1［℃］及びＡｃ3［℃］はＣ、Ｍｎ、Ｓｉの含有量［質量％］によって下記（式２）及び（式３）式から求めたＡｃ1変態温度及びＡｃ3変態温度である。
Ａｃ1＝７６１．３＋２１２Ｃ−４５．８Ｍｎ＋１６．７Ｓｉ ・・・（式２）
Ａｃ3＝９１５−３２５．９Ｃ−３５．９Ｍｎ＋３１．４Ｓｉ ・・・（式３）
(6) A steel slab having the chemical component according to any one of (1) to (3) above is hot-rolled, pickled, and cold-rolled, and then the steel sheet is (Ac 1 [ [Ac] [−100 ° C.] to Ac 1 [° C.] at a rate of temperature increase of 5 ° C./s or more, Ac 1 [° C.] to {Ac 1 [° C.] + 2/3 × (Ac 3 [° C.] − Ac 1 [ ° C])}, the steel sheet is annealed at a residence time of 10 to 300 s where the temperature of the steel sheet is within the temperature range, and is cooled at an average cooling rate of 40 ° C./s or less. A method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance.
Here, Ac 1 [° C.] and Ac 3 [° C.] are the Ac 1 transformation temperature and Ac 3 transformation obtained from the following formulas (Formula 2) and (Formula 3) according to the content [mass%] of C, Mn, and Si, respectively. Temperature.
Ac 1 = 761.3 + 212C-45.8Mn + 16.7Si (Formula 2)
Ac 3 = 915-325.9C-35.9Mn + 31.4Si (Formula 3)
（７） 上記（６）記載の冷却後、３５０〜５００℃の温度範囲で３０ｓ以上保持する過時効処理を施すことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
（８） 上記（６）記載の冷却後、溶融Ｚｎめっきを施すことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
（９） 焼鈍、冷却及び溶融Ｚｎめっきを連続ラインによって施すことを特徴とする上記（８）に記載の加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
（１０） 上記（８）又は（９）記載の溶融Ｚｎめっきを施した後に４５０〜６００℃の温度範囲で１０ｓ以上の熱処理を行うことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
（１１） 上記（６）〜（１０）の何れか１項に記載の方法により製造した冷延鋼板に０．１〜５．０％のスキンパス圧延を施すことを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板の製造方法。
(7) Manufacture of a high strength cold-rolled steel sheet excellent in workability and impact resistance characteristics, characterized by performing an overaging treatment for 30 seconds or more in the temperature range of 350 to 500 ° C. after cooling described in (6) above. Method.
(8) A method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance characteristics, characterized by performing hot-dip Zn plating after the cooling described in (6) above.
(9) The method for producing a high-strength cold-rolled steel sheet having excellent workability and impact resistance as described in (8) above, wherein annealing, cooling and hot-dip Zn plating are performed by a continuous line.
(10) High strength excellent in workability and impact resistance, characterized by performing heat treatment for 10 seconds or more in a temperature range of 450 to 600 ° C. after performing hot-dip Zn plating according to (8) or (9) above A method for producing a cold-rolled steel sheet.
(11) Workability and impact resistance characterized by subjecting a cold-rolled steel sheet produced by the method according to any one of (6) to (10) above to 0.1 to 5.0% skin pass rolling. A method for producing high-strength cold-rolled steel sheets with excellent characteristics.
本発明により、加工性及び耐衝突特性に優れた高強度冷延鋼板の提供が可能になり、特に、生産性を損なわずに安定的に製造できる未再結晶フェライトを積極的に活用した、降伏比が高く、穴拡げ性にも優れた高強度冷延鋼板の提供が可能になり、産業上の貢献が極めて顕著である。 According to the present invention, it is possible to provide a high-strength cold-rolled steel sheet having excellent workability and impact resistance characteristics. It is possible to provide a high-strength cold-rolled steel sheet having a high ratio and excellent hole expansibility, and the industrial contribution is extremely remarkable.
従来、冷延鋼板の金属組織のフェライトの一部を未再結晶フェライトとして残留させるという発想は皆無であった。これは、再結晶が不完全であると冷延鋼板の材質が不均一になると考えられていたためである。 Conventionally, there has been no idea that a part of the ferrite of the metal structure of the cold-rolled steel sheet remains as non-recrystallized ferrite. This is because the material of the cold-rolled steel sheet is considered to be non-uniform when recrystallization is incomplete.
したがって、従来の未再結晶フェライトと硬質第２相からなる冷延鋼板は、未再結晶フェライトの外に焼鈍の加熱時に再結晶したフェライト（再結晶フェライトという。）や焼鈍後の冷却時にオーステナイトから変態したフェライト（変態フェライトという。）が混在したものではなく、フェライトは均質な未再結晶フェライトのみであると考えられる。 Therefore, the conventional cold-rolled steel sheet composed of non-recrystallized ferrite and a hard second phase is composed of ferrite recrystallized during annealing heating (called recrystallized ferrite) in addition to non-recrystallized ferrite and austenite during cooling after annealing. Transformed ferrite (called transformed ferrite) is not mixed, and it is considered that the ferrite is only homogeneous unrecrystallized ferrite.
また、従来、焼鈍の昇温速度を速くし、鋼板の結晶粒径を微細化する製造方法が提案されているが、この方法は、α＋γ二相域での保持によって未再結晶フェライトを完全にオーステナイトに変態させるものであると考えられる。即ち、この従来技術は、焼鈍により未再結晶フェライトを完全にオーステナイトに変態させた後、冷却時にオーステナイトから再変態したフェライトと硬質第２相からなるＤＰ鋼を、未再結晶フェライトを残留させることなく得るものであると推定される。 Conventionally, a manufacturing method has been proposed in which the temperature rise rate of annealing is increased and the crystal grain size of the steel sheet is reduced, but this method completely eliminates unrecrystallized ferrite by holding in the α + γ two-phase region. It is thought to transform to austenite. That is, in this prior art, after non-recrystallized ferrite is completely transformed to austenite by annealing, DP steel composed of ferrite re-transformed from austenite and a hard second phase during cooling remains unrecrystallized ferrite. It is estimated that it can be obtained without.
しかし、焼鈍後の冷却時にオーステナイトをフェライトに変態させると、オーステナイトはフェライトとセメンタイトに分解する。そのため、ベイナイト、マルテンサイト、残留オーステナイトからなる硬質第２相と、セメンタイトを含むフェライトからなるＤＰ鋼となる。そのため、焼鈍時の昇温速度を速くして得られた従来のＤＰ鋼は、局部延性の低下がセメンタイトによって更に助長されていると考えられる。 However, when austenite is transformed into ferrite during cooling after annealing, austenite decomposes into ferrite and cementite. Therefore, it becomes DP steel which consists of a hard 2nd phase which consists of bainite, a martensite, a retained austenite, and a ferrite containing cementite. Therefore, it is considered that the conventional DP steel obtained by increasing the rate of temperature increase during annealing is further promoted by cementite in reducing the local ductility.
一方、本発明は、Ｎｂ、Ｔｉの添加量と焼鈍条件の最適化により、未再結晶フェライトを活用し、延性や穴拡げ性等の加工性を劣化させることなく、鋼板の強度を高めるものである。未再結晶フェライトは、冷間圧延によって圧延方向に延伸されたフェライトの結晶粒が再結晶せず、粒内の転位が回復したものである。そのため、図１に模式的に示したように、未再結晶フェライト１の粒内には転位の回復によって形成されたサブグレイン２を有することが多い。また、未再結晶フェライトの粒内では、冷間圧延による塑性変形のため結晶方位が連続的に変化している。一方、再結晶フェライト及び変態フェライトは、再結晶又は変態によって、粒内の結晶方位はほぼ均一となり、隣接する結晶粒同士の結晶方位は大きく異なっている。
On the other hand, the present invention utilizes non-recrystallized ferrite by optimizing the addition amount of Nb and Ti and annealing conditions, and increases the strength of the steel sheet without degrading workability such as ductility and hole expansibility. is there. Non-recrystallized ferrite is one in which the crystal grains of ferrite stretched in the rolling direction by cold rolling are not recrystallized, and dislocations in the grains are recovered. Therefore, as schematically shown in FIG. 1, the grains of the
更に、本発明では、金属組織をフェライトとパーライトからなるものとし、鋼板の降伏強度と引張強度との比、すなわち、降伏比を上昇させ、耐衝突特性を高めている。そのため、焼入れ性の指標である炭素当量Ｃｅｑを制限して、ベイナイト、マルテンサイトの生成を制限している。 Furthermore, in the present invention, the metal structure is made of ferrite and pearlite, and the ratio between the yield strength and the tensile strength of the steel sheet, that is, the yield ratio is increased to improve the collision resistance. Therefore, the carbon equivalent Ceq, which is an index of hardenability, is limited to limit the generation of bainite and martensite.
したがって、鋼板の成分組成及び焼鈍の条件は本発明において極めて重要であり、再結晶フェライトの細粒化及び未再結晶フェライトを確保するためには、特にＡｃ1以下での昇温速度、最高到達温度及びＡｃ1以上での保持時間、加熱後の冷却速度を制限する必要がある。 Therefore, the composition of the steel sheet and the annealing conditions are extremely important in the present invention, and in order to ensure the recrystallized ferrite fine grain and non-recrystallized ferrite, the rate of temperature rise and the maximum attainment of Ac 1 or less are particularly high. It is necessary to limit the temperature, the holding time at Ac 1 or higher, and the cooling rate after heating.
焼鈍における（Ａｃ1［℃］−１００℃）からＡｃ1［℃］までの昇温速度は５℃／ｓ以上とする。昇温速度を５℃／ｓ以上とする温度の下限を（Ａｃ1［℃］−１００℃）以上としたのは、本発明の鋼の再結晶温度の下限が成分の含有量によって上昇しており、低くとも（Ａｃ1［℃］−１００℃）以上になるためである。また、昇温速度を５℃／ｓ以上とする温度の上限をＡｃ1［℃］としたのは、Ａｃ1［℃］以上の温度ではα−γ変態を生じて、再結晶がほぼ停止するためである。 The temperature increase rate from (Ac 1 [° C.]-100 ° C.) to Ac 1 [° C.] in annealing is 5 ° C./s or more. The lower limit of the temperature at which the rate of temperature rise is 5 ° C./s or higher is set to (Ac 1 [° C.] − 100 ° C.) or higher because the lower limit of the recrystallization temperature of the steel of the present invention is increased by the content of components. This is because at least (Ac 1 [° C.] − 100 ° C.) or higher. Further, the upper limit of the temperature of the heating rate and 5 ° C. / s or more was Ac 1 [° C.] is the Ac 1 [° C.] or higher temperatures occurs the alpha-gamma transformation, recrystallization is substantially stopped Because.
この温度範囲での昇温速度が５℃／ｓ未満の場合、再結晶が十分に進行し、未再結晶フェライトの面積率が著しく減少する。また、昇温速度の上限は設けないが、昇温速度を２０℃／ｓ超としても、再結晶抑制効果が飽和してしまうので、上限を２０℃／ｓとすることが好ましい。 When the rate of temperature increase in this temperature range is less than 5 ° C./s, recrystallization proceeds sufficiently and the area ratio of unrecrystallized ferrite is significantly reduced. Moreover, although the upper limit of a temperature increase rate is not provided, since the recrystallization inhibitory effect will be saturated even if a temperature increase rate exceeds 20 degree-C / s, it is preferable to make an upper limit into 20 degree-C / s.
更に、焼鈍における最高到達温度の下限はＡｃ1［℃］以上とし、上限は、｛Ａｃ1［℃］＋２／３×（Ａｃ3［℃］−Ａｃ1［℃］）｝とする。最高到達温度がＡｃ1未満の場合、フェライトからオーステナイトに変態しないため、パーライトの生成が抑制され、降伏強度の上昇が不十分になる。一方、最高到達温度が｛Ａｃ1［℃］＋２／３×（Ａｃ3［℃］−Ａｃ1［℃］）｝超になると、オーステナイト変態が進行しすぎるため、未再結晶フェライトの確保が困難になり、パーライトが増加して、延性を損なうことがある。 Further, the lower limit of the maximum temperature achieved in annealing is set to Ac 1 [° C.] or more, and the upper limit is set to {Ac 1 [° C.] + 2/3 × (Ac 3 [° C.] − Ac 1 [° C.])}. When the maximum temperature reached is less than Ac 1 , ferrite is not transformed into austenite, so that the formation of pearlite is suppressed and the yield strength is not sufficiently increased. On the other hand, when the maximum temperature reaches {Ac 1 [° C.] + 2/3 × (Ac 3 [° C.] − Ac 1 [° C.])}, the austenite transformation proceeds too much, so it is difficult to secure unrecrystallized ferrite. The pearlite may increase and the ductility may be impaired.
また、鋼板の温度がＡｃ1［℃］以上である温度範囲での滞留時間は１０〜３００ｓとする。これは、以下の理由による。即ち、鋼板の温度がＡｃ1［℃］以上になる時間が１０ｓ未満であると、α−γ変態が十分に進行しないため、パーライトを確保できず、降伏強度が低下する。一方、Ａｃ1［℃］以上での滞留時間が３００ｓを超えると、オーステナイト変態が進行しすぎるため、未再結晶フェライトの確保が困難になり、パーライトが増加して、延性を損なうことがある。 Also, the residence time in the temperature range the temperature of the steel sheet is Ac 1 [° C.] or higher and 10～300S. This is due to the following reason. That is, when the time for which the temperature of the steel sheet is equal to or higher than Ac 1 [° C.] is less than 10 s, the α-γ transformation does not proceed sufficiently, so that pearlite cannot be secured and the yield strength is reduced. On the other hand, if the residence time at Ac 1 [° C.] or more exceeds 300 s, the austenite transformation proceeds too much, so that it becomes difficult to secure unrecrystallized ferrite, pearlite increases, and ductility may be impaired.
なお、Ａｃ1［℃］及びＡｃ3［℃］は、それぞれＡｃ1変態点及びＡｃ3変態点であり、質量％で表されるＣ、Ｍｎ、Ｓｉの含有量により、下記（式２）及び（式３）から求めた温度である。 Ac 1 [° C.] and Ac 3 [° C.] are the Ac 1 transformation point and the Ac 3 transformation point, respectively. Depending on the contents of C, Mn, and Si expressed in mass%, the following (formula 2) and This is the temperature obtained from (Equation 3).
Ａｃ1＝７６１．３＋２１２Ｃ−４５．８Ｍｎ＋１６．７Ｓｉ ・・・（式２）
Ａｃ3＝９１５−３２５．９Ｃ−３５．９Ｍｎ＋３１．４Ｓｉ ・・・（式３）
Ac 1 = 761.3 + 212C-45.8Mn + 16.7Si (Formula 2)
Ac 3 = 915-325.9C-35.9Mn + 31.4Si (Formula 3)
また、本発明では、加工性を劣化させることなく、降伏比を高くするため、降伏強度よりも引張強度の向上への寄与が大きいベイナイト、マルテンサイトの生成を制限している。これにより、降伏比は高くなり、引張強度は若干低下する。したがって、パーライトを活用すると、引張強度の上昇が抑制され、降伏比が高くなるため、加工性の劣化は抑制される。更に、未再結晶フェライトは、粒内に転位の回復によって形成されたサブグレインを有しており、未再結晶フェライトの面積率を増加させると、降伏強度が高くなり、降伏比を高めることができる。 Moreover, in this invention, in order to make a yield ratio high, without reducing workability, the production | generation of the bainite and martensite which have a larger contribution to the improvement of tensile strength than yield strength is restrict | limited. As a result, the yield ratio increases and the tensile strength slightly decreases. Therefore, when pearlite is used, an increase in tensile strength is suppressed and a yield ratio is increased, so that deterioration of workability is suppressed. Furthermore, non-recrystallized ferrite has subgrains formed by dislocation recovery in the grains, and increasing the area ratio of non-recrystallized ferrite increases the yield strength and increases the yield ratio. it can.
また、降伏比を高めるため、本発明では、焼鈍後の平均冷却速度の上限を、４０℃／ｓ以下とすることが必要である。これにより、オーステナイトからパーライトへの変態が促進され、ベイナイトやマルテンサイトの生成が抑制されて、パーライトを十分に確保することができる。 Moreover, in order to raise a yield ratio, in this invention, it is necessary to make the upper limit of the average cooling rate after annealing into 40 degrees C / s or less. Thereby, the transformation from austenite to pearlite is promoted, the generation of bainite and martensite is suppressed, and pearlite can be sufficiently secured.
以下、本発明の限定理由について順次説明する。
まず、本発明における鋼成分及びミクロ組織について説明する。なお、鋼成分の％は質量％を意味する。
Hereinafter, the reasons for limitation of the present invention will be described sequentially.
First, the steel component and microstructure in the present invention will be described. In addition,% of a steel component means the mass%.
Ｃは、強度の増加及びパーライトの生成に寄与する元素であり、狙いとする強度レベルに応じて適量を添加する。Ｃ量は、０．０５％未満であると、高強度を得るのが困難となるため、下限を０．０５％とする。一方、Ｃ量が０．２５％を超えると、成形性や溶接性の劣化を招くため、０．２５％を上限とする。 C is an element that contributes to increase in strength and generation of pearlite, and an appropriate amount is added according to the target strength level. If the amount of C is less than 0.05%, it is difficult to obtain high strength, so the lower limit is made 0.05%. On the other hand, if the amount of C exceeds 0.25%, deterioration of formability and weldability is caused, so 0.25% is made the upper limit.
ＭｎはＡｃ1及び、α−γ変態が完了してオーステナイト単相となる温度であるＡｃ3変態温度（以下、Ａｃ3ともいう。）を低下させる元素であり、本発明において極めて重要である。即ち、Ｍｎ量が少ないと、焼鈍温度を高くする必要が生じ、変態が促進されて未再結晶フェライトの確保が困難になる。また、Ｍｎは、Ｓｉと同様、固溶強化に寄与する元素として強度を増加させるためにも有効である。これらの観点から、Ｍｎ量の下限を０．５０％とする。一方、Ｍｎ量が２．５０％を超えると焼入れ性が高くなり、ベイナイト、マルテンサイトが生成するため、２．５０％を上限とする。降伏比の低下を抑制するには、上限を２．００％以下とすることが好ましい。 Mn is an element that lowers Ac 1 and the Ac 3 transformation temperature (hereinafter also referred to as Ac 3 ), which is the temperature at which the α-γ transformation is completed and becomes an austenite single phase, and is extremely important in the present invention. That is, when the amount of Mn is small, it is necessary to increase the annealing temperature, the transformation is promoted, and it becomes difficult to secure unrecrystallized ferrite. Mn is also effective for increasing the strength as an element contributing to solid solution strengthening, like Si. From these viewpoints, the lower limit of the amount of Mn is 0.50%. On the other hand, when the amount of Mn exceeds 2.50%, the hardenability increases and bainite and martensite are generated, so 2.50% is made the upper limit. In order to suppress a decrease in yield ratio, the upper limit is preferably made 2.00% or less.
Ｓｉは脱酸元素であり、固溶体強化元素として強度を増加させるためにも有効である。しかし、Ｓｉ量が１．００％を超えるとＡｃ1が高くなり過ぎ、焼鈍温度を高くする必要が生じ、変態が促進されて未再結晶フェライトの確保が困難になるため、上限を１．００％以下とする。また、Ｓｉを０．５０％超添加すると溶融Ｚｎめっきを施す際のめっき密着性の低下及び合金化反応の遅延による生産性の低下という問題が生ずることがある。そのため、Ｓｉ量の上限を０．５０％以下とすることが好ましい。Ｓｉ量の下限は規定しないが、０．０１％未満とするには製造コストが高くなるため、下限を０．０１％とすることが好ましい。 Si is a deoxidizing element and is also effective for increasing the strength as a solid solution strengthening element. However, if the amount of Si exceeds 1.00%, Ac 1 becomes too high, and it is necessary to increase the annealing temperature, and the transformation is promoted to make it difficult to secure unrecrystallized ferrite. % Or less. Further, if Si is added in excess of 0.50%, there may be a problem that the plating adhesion is deteriorated when performing hot dip Zn plating and the productivity is lowered due to the delay of the alloying reaction. Therefore, it is preferable that the upper limit of the Si amount is 0.50% or less. The lower limit of the amount of Si is not specified, but if it is less than 0.01%, the manufacturing cost increases, so the lower limit is preferably made 0.01%.
Ａｌは脱酸元素であり、変態点を著しく高める元素であるため、上限を０．２００％とする。なお、Ｓｉ脱酸を行う場合、Ａｌの意図的な添加は不要であり、Ａｌ量の下限は規定しない。 Since Al is a deoxidizing element and is an element that remarkably increases the transformation point, the upper limit is made 0.200%. In addition, when performing Si deoxidation, intentional addition of Al is unnecessary and the minimum of the amount of Al is not prescribed | regulated.
Ｐは不純物であり、粒界に偏析するため、鋼板の靭性の低下や溶接性の劣化を招く。更に、溶融Ｚｎめっき時に合金化反応が極めて遅くなり、生産性が低下する。これらの観点から、Ｐ量の上限を０．１５０％とする。下限は特に限定しないが、Ｐは安価に強度を高める元素であるため、０．００５％以上のＰ量を含有させても良い。 P is an impurity and segregates at the grain boundary, which causes a reduction in toughness and weldability of the steel sheet. Furthermore, the alloying reaction is extremely slow during hot-dip Zn plating, and productivity is reduced. From these viewpoints, the upper limit of the P content is 0.150%. Although a minimum is not specifically limited, Since P is an element which raises an intensity | strength cheaply, you may contain 0.005% or more of P amount.
Ｓは不純物であり、その含有量が０．０１５０％を超えると、熱間割れを誘発したり、加工性を劣化させるので、上限を０．０１５０％とする。Ｓ量の下限は規定しないが、０．０００１％未満とするには製造コストが高くなる。 S is an impurity, and if its content exceeds 0.0150%, hot cracking is induced or workability is deteriorated, so the upper limit is made 0.0150%. The lower limit of the amount of S is not specified, but if it is less than 0.0001%, the manufacturing cost becomes high.
Ｎは不純物であり、Ｎ量が０．０１００％を超えると、靭性や延性の劣化、鋼片の割れの発生が顕著になる。なお、Ｎは、Ｔｉ、Ｎｂの窒化物を生成し、析出に寄与する元素であるため、０．００１０％以上を含有させても良い。 N is an impurity, and when the amount of N exceeds 0.0100%, the deterioration of toughness and ductility and the occurrence of cracks in the steel slab become remarkable. Note that N is an element that forms nitrides of Ti and Nb and contributes to precipitation, so 0.0010% or more may be contained.
Ｎｂ及びＴｉは冷間圧延後の焼鈍工程において、冷間加工によって歪みを導入されたフェライトの再結晶を抑制することにより、未再結晶フェライトの残留を促進する元素である。このような効果を得るためには、Ｎｂ及びＴｉの一方又は双方を合計で０．０４％以上添加することが好ましい。一方、Ｎｂ及びＴｉの一方又は双方の合計の含有量が０．０８％を超えると、再結晶抑制効果が飽和するだけでなく、合金コストが過剰になってしまうため、上限を０．０８％とすることが好ましい。 Nb and Ti are elements that promote the residual of unrecrystallized ferrite by suppressing recrystallization of ferrite introduced with strain by cold working in the annealing process after cold rolling. In order to obtain such an effect, it is preferable to add 0.04% or more of one or both of Nb and Ti in total. On the other hand, if the total content of one or both of Nb and Ti exceeds 0.08%, not only the recrystallization suppression effect is saturated, but the alloy cost becomes excessive, so the upper limit is 0.08%. It is preferable that
Ｃｅｑは鋼板の焼き入れ性を表す指標である。本発明においてはＣｅｑの値が高くなると、降伏比を低下させるベイナイト、マルテンサイトが生成するため、上限を０．４５とする。Ｃｅｑは、下記（式１）によって求められる。
Ｃｅｑ＝Ｃ＋Ｓｉ／２４＋Ｍｎ／６＋Ｎｉ／４０＋Ｃｒ／５＋Ｍｏ／４ ・・・（式１）
ここで、Ｃ、Ｓｉ、Ｍｎ、Ｎｉ、Ｃｒ、Ｍｏは、各元素の含有量［質量％］である。なお、選択的に添加されるＮｉ、Ｃｒ、Ｍｏの含有量が好ましい下限値に満たない場合は０として、（式１）よりＣｅｑを求める。また、強度を確保するためには、Ｃｅｑを０．２以上にすることが好ましい。
Ceq is an index representing the hardenability of the steel sheet. In the present invention, when the value of Ceq increases, bainite and martensite that lower the yield ratio are generated, so the upper limit is set to 0.45. Ceq is calculated | required by the following (Formula 1).
Ceq = C + Si / 24 + Mn / 6 + Ni / 40 + Cr / 5 + Mo / 4 (Formula 1)
Here, C, Si, Mn, Ni, Cr, and Mo are content [mass%] of each element. In addition, when content of Ni, Cr, and Mo added selectively is less than a preferable lower limit, Ceq is calculated | required from (Formula 1) as 0. Further, in order to ensure the strength, it is preferable to set Ceq to 0.2 or more.
本発明では、更に、Ｍｏ、Ｂ、Ｃｒ、Ｎｉの１種又は２種以上を含有させても良い。 In the present invention, one or more of Mo, B, Cr and Ni may be further contained.
Ｍｏ、Ｂ、Ｃｒ及びＮｉは、いずれも固溶強化により引張強度及び降伏強度を高める元素であり、必要に応じて１種又は２種以上を添加しても良い。強度向上の効果を得るためには、それぞれ、Ｍｏ：０．１％以上、Ｂ：０．０００５％以上、Ｃｒ：０．１０％以上、Ｎｉ：０．１０％以上を下限として添加することが好ましい。一方、過剰な添加は合金コストの増加を招くため、それぞれの上限を、Ｍｏ：１．０％以下、Ｂ：０．０１００％以下、Ｃｒ：１．５０％以下、Ｎｉ：１．５０％以下とすることが好ましい。 Mo, B, Cr and Ni are all elements that increase the tensile strength and the yield strength by solid solution strengthening, and one or more of them may be added as necessary. In order to obtain the effect of improving the strength, each of Mo: 0.1% or more, B: 0.0005% or more, Cr: 0.10% or more, Ni: 0.10% or more may be added as lower limits. preferable. On the other hand, excessive addition leads to an increase in alloy costs. Therefore, the upper limits of Mo are 1.0% or less, B is 0.0100% or less, Cr is 1.50% or less, and Ni is 1.50% or less. It is preferable that
本発明によって得られる鋼板のミクロ組織は、フェライトとパーライトからなり、フェライトは、未再結晶フェライト、再結晶フェライト及び変態フェライトの総称である。なお、光学顕微鏡による組織観察では、再結晶フェライトと変態フェライトとの差異は明確ではなく、両者を区別することは困難である。 The microstructure of the steel sheet obtained by the present invention is composed of ferrite and pearlite, and ferrite is a general term for non-recrystallized ferrite, recrystallized ferrite and transformed ferrite. In addition, in the structure observation with an optical microscope, the difference between recrystallized ferrite and transformed ferrite is not clear, and it is difficult to distinguish them.
パーライトは、降伏強度の上昇に寄与する一方で、過剰に存在すると著しく延性が低下するため、下限を１％、上限を３０％とする。 While pearlite contributes to an increase in yield strength, if it is present excessively, ductility is significantly reduced, so the lower limit is made 1% and the upper limit is made 30%.
ミクロ組織は、圧延方向に平行な板厚断面を観察面として試料を採取し、観察面を研磨、ナイタールエッチし、光学顕微鏡で観察すれば良い。本発明の鋼の金属組織は、フェライトとパーライトからなり、ベイナイト及びマルテンサイトを含まないため、光学顕微鏡によって得られたミクロ組織写真を画像解析することによって、パーライトの面積率を、フェライト以外の相の面積率として求めることができる。 The microstructure may be obtained by taking a sample with the cross section of the plate thickness parallel to the rolling direction as an observation surface, polishing the observation surface, performing nital etching, and observing with an optical microscope. Since the metal structure of the steel of the present invention is composed of ferrite and pearlite and does not contain bainite and martensite, the area ratio of pearlite can be determined by analyzing the microstructure image obtained by an optical microscope to obtain a phase ratio other than ferrite. The area ratio can be obtained.
再結晶フェライトと変態フェライトの一方又は双方の面積率は、２０〜７９％とする。これは、再結晶フェライトと変態フェライトの一方又は双方の面積率が、２０％未満では延性が低下し、７９％を超えると強度が低下するためである。 The area ratio of one or both of the recrystallized ferrite and the transformed ferrite is 20 to 79%. This is because the ductility decreases when the area ratio of one or both of the recrystallized ferrite and the transformed ferrite is less than 20%, and the strength decreases when the area ratio exceeds 79%.
未再結晶フェライトは高強度化に寄与することから、その効果を得るためには２０％以上の未再結晶フェライトを含んでいることが必要である。一方、未再結晶フェライトの面積率が５０％を超えると、著しく延性が低下するため、上限を５０％とする。また、未再結晶フェライトの面積率の増加は降伏比の上昇にも寄与する。そのため、未再結晶フェライトの面積率を２５％以上にすることが好ましい。未再結晶フェライトの面積率を増加させるには、焼鈍の加熱速度を高めるか、最高加熱温度を低下させるか、保持時間を短くすれば良い。 Since non-recrystallized ferrite contributes to high strength, it is necessary to contain 20% or more of non-recrystallized ferrite in order to obtain the effect. On the other hand, if the area ratio of non-recrystallized ferrite exceeds 50%, the ductility is remarkably lowered, so the upper limit is made 50%. An increase in the area ratio of non-recrystallized ferrite also contributes to an increase in yield ratio. Therefore, it is preferable to set the area ratio of non-recrystallized ferrite to 25% or more. In order to increase the area ratio of non-recrystallized ferrite, the heating rate of annealing is increased, the maximum heating temperature is decreased, or the holding time is shortened.
未再結晶フェライトとそれ以外のフェライト、即ち再結晶フェライト及び変態フェライトとは、電子後方散乱解析像（Ｅｌｅｃｔｒｏｎ ｂａｃｋ ｓｃａｔｔｅｒｉｎｇ ｐａｔｔｅｒｎ、ＥＢＳＰという。）の結晶方位測定データをＫｅｒｎｅｌ Ａｖｅｒａｇｅ Ｍｉｓｏｒｉｅｎｔａｔｉｏｎ法（ＫＡＭ法）で解析することにより判別することができる。 Non-recrystallized ferrite and other ferrites, that is, recrystallized ferrite and transformed ferrite, are obtained by analyzing the crystal orientation measurement data of an electron back scattering pattern (EBSP) by the Kernel Average Misoration method (KAM method). It can be determined by analysis.
未再結晶フェライトの粒内には、転位は回復しているものの、冷延時の塑性変形によって生じた結晶方位の連続的な変化が存在する。一方、未再結晶フェライトを除くフェライト粒内の結晶方位変化は極めて小さくなる。これは、再結晶及び変態により、隣接する結晶粒の結晶方位は大きく異なるものの、１つの結晶粒内では結晶方位が変化していないためである。ＫＡＭ法では、隣接したピクセル（測定点）との結晶方位差を定量的に示すことができるので、本発明では隣接測定点との平均結晶方位差が１°以内且つ、平均結晶方位差が２°以上あるピクセル間を粒界と定義した時に、結晶粒径が３μｍ以上である粒を未再結晶フェライト以外のフェライト、即ち再結晶フェライト及び変態フェライトと定義する。 In the grains of unrecrystallized ferrite, although dislocations are recovered, there is a continuous change in crystal orientation caused by plastic deformation during cold rolling. On the other hand, the crystal orientation change in the ferrite grains excluding non-recrystallized ferrite becomes extremely small. This is because the crystal orientation does not change in one crystal grain, although the crystal orientation of adjacent crystal grains varies greatly due to recrystallization and transformation. In the KAM method, the crystal orientation difference between adjacent pixels (measurement points) can be quantitatively shown. Therefore, in the present invention, the average crystal orientation difference between adjacent measurement points is within 1 ° and the average crystal orientation difference is 2 When a pixel boundary is defined as a grain boundary, a grain having a crystal grain size of 3 μm or more is defined as ferrite other than unrecrystallized ferrite, that is, recrystallized ferrite and transformed ferrite.
ＥＢＳＰ測定は、焼鈍後の試料の平均結晶粒径の１０分の１の測定間隔で、任意の板断面の板厚方向の１／４厚の位置で１００×１００μｍの範囲において行えば良い。このＥＢＳＰ測定の結果、得られた測定点はピクセルとして出力される。ＥＢＳＰの結晶方位測定に供する試料は、機械研磨等によって鋼板を所定の板厚まで減厚し、次いで電解研磨等によって歪みを除去すると同時に、板厚１／４面が測定面となるように作製する。 The EBSP measurement may be performed in a range of 100 × 100 μm at a 1/4 thickness position in the plate thickness direction of any plate cross section at a measurement interval of 1/10 of the average crystal grain size of the sample after annealing. As a result of the EBSP measurement, the measurement points obtained are output as pixels. Samples to be used for EBSP crystal orientation measurement are prepared so that the steel plate is reduced to a predetermined thickness by mechanical polishing, etc., and then the strain is removed by electrolytic polishing, etc., and at the same time, the 1/4 thickness is the measurement surface. To do.
未再結晶フェライトを含むフェライトの総面積率は、パーライトの面積率の残部であるから、ＥＢＳＰの結晶方位測定に使用した試料をナイタールエッチし、該測定を行った視野の光学顕微鏡写真を同一の倍率で撮影し、得られた組織写真を画像解析して求めれば良い。更に、この組織写真とＥＢＳＰの結晶方位測定の結果を対比させることによって、未再結晶フェライト及び未再結晶フェライト以外のフェライト、即ち、再結晶フェライトと変態フェライトの面積率の合計を求めることもできる。 Since the total area ratio of ferrite including unrecrystallized ferrite is the remainder of the area ratio of pearlite, the sample used for measuring the crystal orientation of EBSP was nital etched, and the optical micrograph of the field of view where the measurement was performed is the same It is only necessary to obtain the image of the obtained tissue photograph by image analysis. Furthermore, by comparing the result of the crystal orientation measurement of this structural photograph and EBSP, the total area ratio of non-recrystallized ferrite and ferrite other than non-recrystallized ferrite, that is, recrystallized ferrite and transformed ferrite can be obtained. .
また、耐衝突特性を向上させるには、降伏比を０．７０以上に高めることが好ましい。降伏強度を上昇させ、降伏比を０．７０以上に高めると、耐衝突特性が顕著に向上する。好ましくは０．７５以上、更に好ましくは０．８０以上である。なお、降伏比を高めるには、未再結晶フェライトの面積率を増加させることが好ましい。 In order to improve the collision resistance, it is preferable to increase the yield ratio to 0.70 or more. When the yield strength is increased and the yield ratio is increased to 0.70 or more, the collision resistance is remarkably improved. Preferably it is 0.75 or more, More preferably, it is 0.80 or more. In order to increase the yield ratio, it is preferable to increase the area ratio of non-recrystallized ferrite.
耐食性が要求される用途に適用する場合は、溶融Ｚｎめっきを設けることが好ましい。更に、耐食性の向上が要求される場合は、合金化処理を施した合金化溶融Ｚｎめっきを設けることが好ましい。 When applied to applications requiring corrosion resistance, hot-dip Zn plating is preferably provided. Furthermore, when improvement in corrosion resistance is required, it is preferable to provide alloyed hot-dip Zn plating that has been subjected to alloying treatment.
次に、製造方法及びその好ましい条件について説明する。 Next, a manufacturing method and its preferable conditions are demonstrated.
熱間圧延に供する鋼片は常法で製造すれば良く、鋼を溶製し、鋳造すれば良い。生産性の観点からは、連続鋳造が好ましく、薄スラブキャスター等で製造しても良い。また、鋳造後直ちに熱間圧延を行う連続鋳造―直接圧延のようなプロセスでも良い。熱間圧延は常法で行えば良く、圧延温度、圧下率、冷却速度、巻取温度等の条件は特に規定しない。熱間圧延後、鋼板を冷間圧延、焼鈍し、冷延鋼板とする。 The steel piece to be subjected to hot rolling may be manufactured by a conventional method, and the steel may be melted and cast. From the viewpoint of productivity, continuous casting is preferable, and it may be manufactured with a thin slab caster or the like. Further, a process such as continuous casting-direct rolling in which hot rolling is performed immediately after casting may be used. Hot rolling may be performed by a conventional method, and conditions such as rolling temperature, rolling reduction, cooling rate, and winding temperature are not particularly specified. After hot rolling, the steel sheet is cold-rolled and annealed to obtain a cold-rolled steel sheet.
冷間圧延の圧下率は特に規定しないが、１０％未満の冷間圧延率では、板厚制御が難しく形状不良の原因となるため、その下限を１０％以上とすることが好ましい。一方、冷間圧延率が９０％超になると、圧延ロールへの負荷が大きくなる上、再結晶が促進されて未再結晶フェライトを確保するために、焼鈍の昇温速度を大きくすることが必要になる。そのため、冷間圧延の圧下率の上限は、９０％以下とすることが好ましい。 Although the rolling reduction ratio of cold rolling is not particularly defined, it is preferable to set the lower limit to 10% or more because cold rolling reduction of less than 10% makes it difficult to control the plate thickness and causes shape defects. On the other hand, if the cold rolling rate exceeds 90%, the load on the rolling roll increases, and it is necessary to increase the heating rate of annealing in order to promote recrystallization and secure unrecrystallized ferrite. become. Therefore, the upper limit of the cold rolling reduction is preferably 90% or less.
本発明において、冷間圧延後の焼鈍は極めて重要であり、上述の条件で行うことが必要である。焼鈍は、昇温速度、加熱時間を制御するため、連続焼鈍設備によって行うことが好ましい。また、昇温速度を速くするために、高周波加熱装置、通電加熱装置を併用しても良い。焼鈍において、Ａｃ1以上での滞留時間は、鋼板の温度がＡｃ1以上である時間の合計であり、加熱炉の設定温度と炉の長さ、通板速度によって制御することができる。 In the present invention, annealing after cold rolling is extremely important, and it is necessary to carry out under the above-mentioned conditions. Annealing is preferably performed by continuous annealing equipment in order to control the rate of temperature rise and the heating time. Further, in order to increase the rate of temperature rise, a high-frequency heating device or an electric heating device may be used in combination. In annealing, the residence time in the Ac 1 or more, the sum of the time the temperature of the steel sheet is Ac 1 or more, the length of the set temperature and the furnace of the heating furnace can be controlled by the sheet passing speed.
焼鈍後の冷却速度は、水等、冷媒の吹付け、送風、ミスト等による強制冷却により、適宜制御すれば良い。また、平均冷却速度は、強制冷却の開始から停止までの冷却速度である。強制冷却は、連続焼鈍ラインで焼鈍を行う際には、室温で停止しても良いが、２００〜３５０℃で停止することが多い。なお、焼鈍後、過時効処理を施す場合には、３５０〜５００℃の範囲内で冷却を停止しても良い。連続焼鈍−溶融Ｚｎめっきラインで焼鈍を行う場合は、４００〜５００℃まで冷却することが多い。 The cooling rate after annealing may be appropriately controlled by forced cooling with water or the like, blowing of refrigerant, blowing air, mist, or the like. The average cooling rate is a cooling rate from the start to the stop of forced cooling. The forced cooling may be stopped at room temperature when annealing is performed in a continuous annealing line, but often stops at 200 to 350 ° C. In addition, when performing an overaging process after annealing, you may stop cooling within the range of 350-500 degreeC. In the case of performing annealing in a continuous annealing-hot Zn plating line, it is often cooled to 400 to 500 ° C.
焼鈍後は、３５０〜５００℃に保持する過時効処理を行うことが好ましい。冷却後の鋼板に残留オーステナイトが存在していた場合、３５０〜５００℃の温度範囲で、３０ｓ以上保持する過時効処理を行うと、残留オーステナイトがフェライトとセメンタイトに分解されるため、降伏比を上昇させることができる。 After annealing, it is preferable to perform an overaging treatment that is maintained at 350 to 500 ° C. When retained austenite is present in the steel sheet after cooling, the yield ratio is increased because the retained austenite is decomposed into ferrite and cementite when subjected to an overaging treatment at a temperature range of 350 to 500 ° C. for 30 seconds or more. Can be made.
焼鈍後、必要に応じて、過時効処理、溶融Ｚｎめっき又は合金化溶融Ｚｎめっきを施しても良い。また、Ｚｎめっきの組成は特に限定するものではなく、Ｚｎの他、Ｆｅ、Ａｌ、Ｍｎ、Ｃｒ、Ｍｇ、Ｐｂ、Ｓｎ、Ｎｉ等を必要に応じて添加しても構わない。なお、めっきは、焼鈍と別工程で行っても良いが、生産性の観点から、焼鈍とめっきを連続して行う、連続焼鈍−溶融Ｚｎめっきラインによって行うことが好ましい。この場合も、未再結晶フェライトを確保するためには、焼鈍を上記の条件で行うことが必要である。 After annealing, an overaging treatment, hot dip Zn plating, or alloyed hot dip Zn plating may be applied as necessary. Further, the composition of the Zn plating is not particularly limited, and in addition to Zn, Fe, Al, Mn, Cr, Mg, Pb, Sn, Ni, or the like may be added as necessary. The plating may be performed in a separate process from the annealing, but from the viewpoint of productivity, it is preferable that the plating is performed by a continuous annealing-hot Zn plating line in which annealing and plating are continuously performed. Also in this case, in order to ensure non-recrystallized ferrite, it is necessary to perform annealing under the above conditions.
合金化処理を行う場合は、４５０〜６００℃の温度範囲で行うことが好ましい。これは、４５０℃未満では合金化が十分に進行せず、また、６００℃超では過度に合金化が進行し、めっき層が脆化して、プレス等の加工によってめっきが剥離する等の問題を誘発することがあるためである。合金化処理の時間は、１０ｓ未満では合金化が十分に進行しないことがあるため、１０ｓ以上とすることが好ましい。また、合金化処理の時間の上限は特に規定しないが、生産効率の観点から１００ｓ以内とすることが好ましい。
また、生産性の観点から、連続焼鈍−溶融Ｚｎめっきラインに合金化処理炉を連続して設け、焼鈍、めっき及び合金化処理を連続して行うことが好ましい。
When performing an alloying process, it is preferable to carry out in the temperature range of 450-600 degreeC. This is because the alloying does not proceed sufficiently below 450 ° C., and the alloying proceeds excessively above 600 ° C., the plating layer becomes brittle, and the plating peels off by processing such as pressing. This is because it may trigger. When the alloying treatment time is less than 10 s, alloying may not proceed sufficiently. Further, the upper limit of the alloying time is not particularly defined, but is preferably within 100 s from the viewpoint of production efficiency.
Further, from the viewpoint of productivity, it is preferable to continuously provide an alloying treatment furnace in the continuous annealing-hot Zn plating line and to perform annealing, plating, and alloying treatment continuously.
表１に示す組成を有する鋼を溶製し、鋳造して得られた鋼片を、１２５０℃で再加熱した後、常法に従って熱間圧延を行った。この時、仕上げ温度は９００℃、巻取温度は６００℃とした。その後、６０％の圧下率で冷間圧延を施した後、表２に示す条件で焼鈍を行った。なお、表１の［−］は、成分を意図的に添加していないことを意味する。また、表１には、Ａｃ1［℃］、Ａｃ3［℃］及びＣｅｑの計算値も示した。表２の昇温速度は、（Ａｃ1［℃］−１００℃）からＡｃ1［℃］までの温度の上昇に要した時間によって計算した。表２の冷却速度は、強制冷却を行った温度範囲と所要時間から求めた。 Steel pieces obtained by melting and casting steel having the composition shown in Table 1 were reheated at 1250 ° C. and then hot-rolled according to a conventional method. At this time, the finishing temperature was 900 ° C., and the winding temperature was 600 ° C. Then, after performing cold rolling at a reduction rate of 60%, annealing was performed under the conditions shown in Table 2. In addition, [-] in Table 1 means that no component is intentionally added. Table 1 also shows the calculated values of Ac 1 [° C.], Ac 3 [° C.] and Ceq. The heating rate in Table 2 was calculated according to the time required for the temperature increase from (Ac 1 [° C.]-100 ° C.) to Ac 1 [° C.]. The cooling rate in Table 2 was determined from the temperature range and required time for forced cooling.
表２に示す冷延鋼板のうち、製造Ｎｏ．２及び６については、焼鈍工程後、Ｚｎめっき浴に浸漬後、製造Ｎｏ．６については更に５００℃で２０ｓ間の合金化処理を施した。更に、表２に示す冷延鋼板のうち、製造Ｎｏ．９については、焼鈍後、３００℃まで冷却し、３００℃で４００ｓ保持する過時効処理を行った後、１０℃／ｓで室温まで冷却した。また、製造Ｎｏ．１３は、１．０％のスキンパス圧延を施した例である。 Among the cold-rolled steel sheets shown in Table 2, production No. About No. 2 and 6, after an annealing process, after immersion in Zn plating bath, manufacture No. 6 was further alloyed at 500 ° C. for 20 s. Furthermore, among the cold-rolled steel sheets shown in Table 2, the production No. About No. 9, after annealing, it cooled to 300 degreeC and performed the overaging process which hold | maintains 400 s at 300 degreeC, Then, it cooled to room temperature at 10 degreeC / s. In addition, production No. 13 is an example in which 1.0% skin pass rolling was performed.
製造後の冷延鋼板から、幅方向（ＴＤ方向という。）を長手方向としてＪＩＳ Ｚ ２２０１の５号引張試験片を採取し、ＪＩＳ Ｚ ２２４１に準拠してＴＤ方向の引張特性を評価した。なお、ｔ−Ｅｌ［％］は破断伸びであり、降伏比は、降伏強度を引張強度で除した値である。また、穴拡げ試験は日本鉄鋼連盟規格ＪＦＳ Ｔ １００１−１９９６記載の試験方法に準拠して行ない、穴拡げ率λを評価した。 From the cold-rolled steel sheet after production, a No. 5 tensile test piece of JIS Z 2201 was taken with the width direction (referred to as the TD direction) as the longitudinal direction, and the tensile characteristics in the TD direction were evaluated according to JIS Z 2241. In addition, t-El [%] is elongation at break, and the yield ratio is a value obtained by dividing the yield strength by the tensile strength. Moreover, the hole expansion test was conducted in accordance with the test method described in the Japan Iron and Steel Federation Standard JFS T 1001-1996, and the hole expansion rate λ was evaluated.
ミクロ組織は、圧延方向に平行な板厚断面を観察面として試料を採取し、観察面を研磨、ナイタールエッチ、必要に応じてレペラーエッチ、ピクラールエッチし、光学顕微鏡で観察した。なお、光学顕微鏡による観察の結果、一部の鋼板の金属組織にベイナイト、マルテンサイトが観察された。そのため、得られたミクロ組織写真を画像解析し、パーライト、ベイナイト、マルテンサイトの面積率を求めた。以下では、ベイナイト、マルテンサイトの一方又は双方の面積率の合計を、低温変態相面積率という。未再結晶フェライトの面積率は、ＥＢＳＰによって測定した。 The microstructure was obtained by taking a sample with a cross section of the plate thickness parallel to the rolling direction as an observation surface, polishing the observation surface, nital etching, if necessary, repeller etching, picral etching, and observing with an optical microscope. As a result of observation with an optical microscope, bainite and martensite were observed in the metal structures of some steel sheets. Therefore, the obtained microstructure photograph was subjected to image analysis, and the area ratio of pearlite, bainite, and martensite was obtained. Hereinafter, the total area ratio of one or both of bainite and martensite is referred to as a low-temperature transformation phase area ratio. The area ratio of unrecrystallized ferrite was measured by EBSP.
結果を表３に示す。ここで、表３に示した金属組織の残部は、再結晶フェライトと変態フェライトの一方又は双方の合計の面積率である。表３に示したように、本発明の化学成分を有する鋼を適正な条件で熱延及び冷延し、更に、適切な条件で焼鈍することにより、更に、過時効処理、Ｚｎめっき、合金化処理、スキンパス圧延を施しても加工性及び耐衝突特性に優れた高強度冷延鋼板を得ることが可能である。 The results are shown in Table 3. Here, the balance of the metal structure shown in Table 3 is the total area ratio of one or both of the recrystallized ferrite and the transformed ferrite. As shown in Table 3, the steel having the chemical components of the present invention is hot-rolled and cold-rolled under appropriate conditions, and further annealed under appropriate conditions, so that overaging treatment, Zn plating, and alloying are further performed. It is possible to obtain a high-strength cold-rolled steel sheet that is excellent in workability and impact resistance characteristics even when subjected to treatment and skin pass rolling.
一方、鋼Ｎｏ．ＧはＣ量が少ないため、強度が低下している。また、鋼Ｎｏ．ＨはＣｅｑが高いため、焼き入れ性が高くなりすぎて、低温変態相面積率が増加し、降伏比及び穴拡げ率λが低下している。鋼Ｎｏ．ＩはＮｂ及びＴｉ量が少ないため、未再結晶フェライトが少なくなり、強度及び降伏比が低下し、穴拡げ率λが低下している。 On the other hand, Steel No. Since G has a small amount of C, its strength is reduced. Steel No. Since H has a high Ceq, the hardenability becomes too high, the low-temperature transformation phase area ratio increases, and the yield ratio and the hole expansion ratio λ decrease. Steel No. Since I has a small amount of Nb and Ti, unrecrystallized ferrite decreases, strength and yield ratio decrease, and hole expansion ratio λ decreases.
また、製造Ｎｏ．３は、（Ａｃ1［℃］−１００℃）からＡｃ1［℃］までの昇温速度が遅く、未再結晶フェライトが少なくなり、強度が低下している。製造Ｎｏ．４は、焼鈍の最高到達温度が高く、製造Ｎｏ．７は、Ａｃ1［℃］以上での滞留時間が長いため、未再結晶フェライトが少なく、強度が低下している。また、これらは、パーライトが増加しているため、同等の強度を有する鋼板と比較すると、延性がやや低下している。 In addition, production No. In No. 3, the rate of temperature increase from (Ac 1 [° C.]-100 ° C.) to Ac 1 [° C.] is slow, the amount of unrecrystallized ferrite is reduced, and the strength is reduced. Production No. No. 4 has a high maximum temperature for annealing. No. 7 has a long residence time at Ac 1 [° C.] or higher, so that there is little unrecrystallized ferrite and the strength is lowered. Moreover, since pearlite is increasing in these, ductility has fallen a little compared with the steel plate which has equivalent intensity | strength.
製造Ｎｏ．１２は、焼鈍の最高到達温度からの平均冷却速度が速いため、パーライトの割合が少なく、低温変態相面積率が高くなり、高強度ではあるものの、降伏比が低下し、穴拡げ率λが低下している。 Production No. No. 12, because the average cooling rate from the highest annealing temperature is high, the ratio of pearlite is small, the low-temperature transformation phase area ratio is high, and the strength is high, but the yield ratio is decreased and the hole expansion ratio λ is decreased. is doing.
製造Ｎｏ．１５は、焼鈍の最高到達温度が低く、製造Ｎｏ．１８は、Ａｃ1［℃］以上での滞留時間が短く、パーライトが十分に得られなかったため、降伏比が低下している。 Production No. No. 15 has a low maximum temperature for annealing. No. 18 has a short residence time at Ac 1 [° C.] or more, and pearlite was not sufficiently obtained, so that the yield ratio was lowered.
１ 未再結晶フェライト
２ サブグレイン
1
Claims (11)
Ｃ ：０．０５〜０．２５％、
Ｍｎ：０．５０〜２．５０％、
を含有し、
Ｓｉ：１．００％以下、
Ａｌ：０．２００％以下、
Ｐ ：０．１５０％以下、
Ｓ ：０．０１５０％以下、
Ｎ ：０．０１００％以下
に制限し、更に、
Ｎｂ、Ｔｉの一方又は双方を合計で０．０４〜０．０８％
含有し、下記（式１）によって求められるＣｅｑが０．４５以下であり、残部が鉄及び不可避的不純物からなり、金属組織がフェライトとパーライトからなり、前記フェライトが再結晶フェライト、変態フェライトの一方又は双方と未再結晶フェライトからなり、前記未再結晶フェライトの面積率が２０〜５０％であり、前記再結晶フェライト、前記変態フェライトの一方又は双方の面積率が２０〜７９％であり、前記パーライトの面積率が１〜３０％であることを特徴とする加工性及び耐衝突特性に優れた高強度冷延鋼板。
Ｃｅｑ＝Ｃ＋Ｓｉ／２４＋Ｍｎ／６＋Ｎｉ／４０＋Ｃｒ／５＋Ｍｏ／４ ・・・（式１）
ここで、Ｃ、Ｓｉ、Ｍｎ、Ｎｉ、Ｃｒ、Ｍｏは、各元素の含有量［質量％］である。 % By mass
C: 0.05 to 0.25%,
Mn: 0.50 to 2.50%,
Containing
Si: 1.00% or less,
Al: 0.200% or less,
P: 0.150% or less,
S: 0.0150% or less,
N: limited to 0.0100% or less, and
One or both of Nb and Ti in total 0.04 to 0.08%
Ceq obtained by the following (formula 1) is 0.45 or less, the balance is made of iron and inevitable impurities, the metal structure is made of ferrite and pearlite, and the ferrite is one of recrystallized ferrite and transformed ferrite. Or both and non-recrystallized ferrite, the area ratio of the non-recrystallized ferrite is 20-50%, the area ratio of one or both of the recrystallized ferrite and the transformation ferrite is 20-79%, A high-strength cold-rolled steel sheet excellent in workability and impact resistance, characterized in that the area ratio of pearlite is 1 to 30%.
Ceq = C + Si / 24 + Mn / 6 + Ni / 40 + Cr / 5 + Mo / 4 (Formula 1)
Here, C, Si, Mn, Ni, Cr, and Mo are content [mass%] of each element.
Ｍｏ：０．１〜１．０％、
Ｂ ：０．０００５〜０．０１００％、
Ｃｒ：０．１０〜１．５０％、
Ｎｉ：０．１０〜１．５０％
のうち、１種又は２種以上を含有することを特徴とする請求項１又は２の何れか１項に記載の加工性及び耐衝突特性に優れた高強度冷延鋼板。 % By mass
Mo: 0.1 to 1.0%,
B: 0.0005 to 0.0100%,
Cr: 0.10 to 1.50%,
Ni: 0.10 to 1.50%
The high-strength cold-rolled steel sheet excellent in workability and impact resistance characteristics according to any one of claims 1 and 2, wherein one or more of them are contained.
ここで、Ａｃ1［℃］及びＡｃ3［℃］はＣ、Ｍｎ、Ｓｉの含有量［質量％］によって下記（式２）及び（式３）式から求めたＡｃ1変態温度及びＡｃ3変態温度である。
Ａｃ1＝７６１．３＋２１２Ｃ−４５．８Ｍｎ＋１６．７Ｓｉ ・・・（式２）
Ａｃ3＝９１５−３２５．９Ｃ−３５．９Ｍｎ＋３１．４Ｓｉ ・・・（式３） A steel slab having the chemical composition according to any one of claims 1 to 3 is hot-rolled, pickled, and cold-rolled, and then the steel plate is (Ac 1 [° C] -100 ° C). temperature of Ac 1 [℃] Ac 1 [ ℃] a heating rate of up to a 5 ° C. / s or higher ~ {Ac 1 [℃] + 2/3 × (Ac 3 [℃] -Ac 1 [℃])} from Workability and collision resistance characterized by heating within a range, annealing with the residence time of the steel sheet being within the temperature range as 10 to 300 s, and cooling with an average cooling rate of 40 ° C./s or less A method for producing high-strength cold-rolled steel sheets with excellent characteristics.
Here, Ac 1 [° C.] and Ac 3 [° C.] are the Ac 1 transformation temperature and Ac 3 transformation obtained from the following formulas (Formula 2) and (Formula 3) according to the content [mass%] of C, Mn, and Si, respectively. Temperature.
Ac 1 = 761.3 + 212C-45.8Mn + 16.7Si (Formula 2)
Ac 3 = 915-325.9C-35.9Mn + 31.4Si (Formula 3)
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2008027856A JP4995109B2 (en) | 2008-02-07 | 2008-02-07 | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same |
Applications Claiming Priority (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2008027856A JP4995109B2 (en) | 2008-02-07 | 2008-02-07 | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2009185355A true JP2009185355A (en) | 2009-08-20 |
| JP4995109B2 JP4995109B2 (en) | 2012-08-08 |
Family
ID=41068880
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2008027856A Active JP4995109B2 (en) | 2008-02-07 | 2008-02-07 | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP4995109B2 (en) |
Cited By (19)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| WO2012053642A1 (en) * | 2010-10-22 | 2012-04-26 | 新日本製鐵株式会社 | Method for manufacturing hot stamped body having vertical wall, and hot stamped body having vertical wall |
| WO2012053636A1 (en) * | 2010-10-22 | 2012-04-26 | 新日本製鐵株式会社 | Process for producing hot stamp molded article, and hot stamp molded article |
| EP2500445A1 (en) * | 2009-11-09 | 2012-09-19 | Nippon Steel Corporation | High-strength steel sheet having excellent processability and paint bake hardenability, and method for producing of high-strength steel sheet |
| KR101235535B1 (en) | 2011-06-16 | 2013-02-28 | 현대하이스코 주식회사 | High strength steel sheet with low yield ratio and method of manufacturing the steel sheet |
| WO2013039272A1 (en) * | 2011-09-16 | 2013-03-21 | 주식회사 포스코 | Thin cold-rolled steel plate having high strength and high formability, and preparation method thereof |
| WO2013088666A1 (en) | 2011-12-12 | 2013-06-20 | Ｊｆｅスチール株式会社 | High-yield-ratio high-strength cold-rolled steel sheet and method for producing same |
| JP2013151708A (en) * | 2010-12-27 | 2013-08-08 | Nippon Steel & Sumitomo Metal Corp | Method for manufacturing hot stamped body having small variation in hardness, and the hot stamped body |
| JP2013204141A (en) * | 2012-03-29 | 2013-10-07 | Jfe Steel Corp | High-strength steel sheet excellent in surface quality and manufacturing method thereof |
| KR20150023745A (en) * | 2012-06-27 | 2015-03-05 | 제이에프이 스틸 가부시키가이샤 | Steel sheet for soft nitriding and process for producing same |
| CN105274432A (en) * | 2014-06-11 | 2016-01-27 | 鞍钢股份有限公司 | 600 MPa grade high-yield-ratio high-plasticity cold-rolled steel plate and manufacturing method thereof |
| WO2017051998A1 (en) * | 2015-09-22 | 2017-03-30 | 현대제철 주식회사 | Plated steel plate and manufacturing method thereof |
| US9896736B2 (en) | 2010-10-22 | 2018-02-20 | Nippon Steel & Sumitomo Metal Corporation | Method for manufacturing hot stamped body having vertical wall and hot stamped body having vertical wall |
| KR20180031738A (en) | 2015-08-31 | 2018-03-28 | 신닛테츠스미킨 카부시키카이샤 | Steel plate |
| KR20180038030A (en) | 2015-08-21 | 2018-04-13 | 신닛테츠스미킨 카부시키카이샤 | Steel plate |
| WO2018147211A1 (en) | 2017-02-13 | 2018-08-16 | Ｊｆｅスチール株式会社 | Cold rolled steel sheet and method for manufacturing same |
| WO2019093689A1 (en) * | 2017-11-07 | 2019-05-16 | 주식회사 포스코 | High-strength, low-toughness cold-rolled steel sheet having excellent fracture properties, and method for producing same |
| WO2021149676A1 (en) | 2020-01-22 | 2021-07-29 | 日本製鉄株式会社 | Steel sheet and method for producing same |
| CN114207172A (en) * | 2019-07-31 | 2022-03-18 | 杰富意钢铁株式会社 | High-strength steel sheet, high-strength member, and method for producing same |
| WO2023052814A1 (en) * | 2021-09-29 | 2023-04-06 | Arcelormittal | Cold rolled and heat treated steel sheet and a method of manufacturing thereof |
Families Citing this family (5)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| KR101989371B1 (en) | 2015-03-25 | 2019-06-14 | 제이에프이 스틸 가부시키가이샤 | High-strength steel sheet and method for producing the same |
| EP3255167B1 (en) | 2015-03-25 | 2019-09-25 | JFE Steel Corporation | High-strength steel sheet and method for producing the same |
| EP3305932B1 (en) | 2015-05-29 | 2020-02-12 | JFE Steel Corporation | High strength steel sheet and method for producing same |
| WO2017029815A1 (en) | 2015-08-19 | 2017-02-23 | Ｊｆｅスチール株式会社 | High-strength steel sheet and production method for same |
| WO2017029814A1 (en) | 2015-08-19 | 2017-02-23 | Ｊｆｅスチール株式会社 | High-strength steel sheet and production method for same |
Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2004131771A (en) * | 2002-10-09 | 2004-04-30 | Nippon Steel Corp | Method for manufacturing cold rolled steel sheet having excellent shape fixability |
| JP2006152362A (en) * | 2004-11-29 | 2006-06-15 | Jfe Steel Kk | High-strength steel plate having excellent rigidity and production method thereof |
| JP2007284783A (en) * | 2006-03-20 | 2007-11-01 | Nippon Steel Corp | High strength cold rolled steel sheet and its production method |
| JP2008190032A (en) * | 2007-01-10 | 2008-08-21 | Nippon Steel Corp | High-strength cold-rolled steel sheet excellent in workability and crash resistance and its production method |
-
2008
- 2008-02-07 JP JP2008027856A patent/JP4995109B2/en active Active
Patent Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2004131771A (en) * | 2002-10-09 | 2004-04-30 | Nippon Steel Corp | Method for manufacturing cold rolled steel sheet having excellent shape fixability |
| JP2006152362A (en) * | 2004-11-29 | 2006-06-15 | Jfe Steel Kk | High-strength steel plate having excellent rigidity and production method thereof |
| JP2007284783A (en) * | 2006-03-20 | 2007-11-01 | Nippon Steel Corp | High strength cold rolled steel sheet and its production method |
| JP2008190032A (en) * | 2007-01-10 | 2008-08-21 | Nippon Steel Corp | High-strength cold-rolled steel sheet excellent in workability and crash resistance and its production method |
Cited By (44)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| EP2500445A1 (en) * | 2009-11-09 | 2012-09-19 | Nippon Steel Corporation | High-strength steel sheet having excellent processability and paint bake hardenability, and method for producing of high-strength steel sheet |
| EP2500445A4 (en) * | 2009-11-09 | 2017-05-17 | Nippon Steel & Sumitomo Metal Corporation | High-strength steel sheet having excellent processability and paint bake hardenability, and method for producing of high-strength steel sheet |
| KR101513378B1 (en) | 2010-10-22 | 2015-04-17 | 신닛테츠스미킨 카부시키카이샤 | Steel sheet for hot stamping and method for producing steel sheet for hot stamping |
| WO2012053637A1 (en) * | 2010-10-22 | 2012-04-26 | 新日本製鐵株式会社 | Steel sheet and steel sheet production process |
| JP5293902B2 (en) * | 2010-10-22 | 2013-09-18 | 新日鐵住金株式会社 | Hot stamping steel plate and method for manufacturing hot stamping steel plate |
| US9598745B2 (en) | 2010-10-22 | 2017-03-21 | Nippon Steel & Sumitomo Metal Corporation | Method for manufacturing hot stamped body and hot stamped body |
| WO2012053642A1 (en) * | 2010-10-22 | 2012-04-26 | 新日本製鐵株式会社 | Method for manufacturing hot stamped body having vertical wall, and hot stamped body having vertical wall |
| JP5224010B2 (en) * | 2010-10-22 | 2013-07-03 | 新日鐵住金株式会社 | Method for producing hot stamping molded body having vertical wall and hot stamping molded body having vertical wall |
| US9512499B2 (en) | 2010-10-22 | 2016-12-06 | Nippon Steel & Sumitomo Metal Corporation | Method for manufacturing hot stamped body having vertical wall and hot stamped body having vertical wall |
| CN103261452A (en) * | 2010-10-22 | 2013-08-21 | 新日铁住金株式会社 | Method for manufacturing hot stamped body having vertical wall, and hot stamped body having vertical wall |
| US20130220490A1 (en) * | 2010-10-22 | 2013-08-29 | Kunio Hayashi | Steel sheet and method for manufacturing steel sheet |
| JP5547287B2 (en) * | 2010-10-22 | 2014-07-09 | 新日鐵住金株式会社 | Manufacturing method of hot stamping molded body and hot stamping molded body |
| US9896736B2 (en) | 2010-10-22 | 2018-02-20 | Nippon Steel & Sumitomo Metal Corporation | Method for manufacturing hot stamped body having vertical wall and hot stamped body having vertical wall |
| US9840751B2 (en) | 2010-10-22 | 2017-12-12 | Nippon Steel & Sumitomo Metal Corporation | Method for manufacturing hot stamped body and hot stamped body |
| CN103314120A (en) * | 2010-10-22 | 2013-09-18 | 新日铁住金株式会社 | Process for producing hot stamp molded article, and hot stamp molded article |
| CN103314120B (en) * | 2010-10-22 | 2014-11-05 | 新日铁住金株式会社 | Process for producing hot stamp molded article, and hot stamp molded article |
| WO2012053636A1 (en) * | 2010-10-22 | 2012-04-26 | 新日本製鐵株式会社 | Process for producing hot stamp molded article, and hot stamp molded article |
| US10030280B2 (en) | 2010-10-22 | 2018-07-24 | Nippon Steel & Sumitomo Metal Corporation | Steel sheet and method for manufacturing steel sheet |
| JP2013151708A (en) * | 2010-12-27 | 2013-08-08 | Nippon Steel & Sumitomo Metal Corp | Method for manufacturing hot stamped body having small variation in hardness, and the hot stamped body |
| KR101235535B1 (en) | 2011-06-16 | 2013-02-28 | 현대하이스코 주식회사 | High strength steel sheet with low yield ratio and method of manufacturing the steel sheet |
| JP2014530296A (en) * | 2011-09-16 | 2014-11-17 | ポスコ | Ultra-thin cold-rolled steel sheet having high strength and high formability and manufacturing method thereof |
| WO2013039272A1 (en) * | 2011-09-16 | 2013-03-21 | 주식회사 포스코 | Thin cold-rolled steel plate having high strength and high formability, and preparation method thereof |
| US9994941B2 (en) | 2011-12-12 | 2018-06-12 | Jfe Steel Corporation | High strength cold rolled steel sheet with high yield ratio and method for producing the same |
| WO2013088666A1 (en) | 2011-12-12 | 2013-06-20 | Ｊｆｅスチール株式会社 | High-yield-ratio high-strength cold-rolled steel sheet and method for producing same |
| JP2013204141A (en) * | 2012-03-29 | 2013-10-07 | Jfe Steel Corp | High-strength steel sheet excellent in surface quality and manufacturing method thereof |
| KR101701652B1 (en) | 2012-06-27 | 2017-02-01 | 제이에프이 스틸 가부시키가이샤 | Steel sheet for soft-nitriding and method for manufacturing the same |
| EP2868764A4 (en) * | 2012-06-27 | 2016-04-06 | Jfe Steel Corp | Steel sheet for soft nitriding and process for producing same |
| KR20150023745A (en) * | 2012-06-27 | 2015-03-05 | 제이에프이 스틸 가부시키가이샤 | Steel sheet for soft nitriding and process for producing same |
| US10077489B2 (en) | 2012-06-27 | 2018-09-18 | Jfe Steel Corporation | Steel sheet for soft-nitriding and method for manufacturing the same |
| CN105274432A (en) * | 2014-06-11 | 2016-01-27 | 鞍钢股份有限公司 | 600 MPa grade high-yield-ratio high-plasticity cold-rolled steel plate and manufacturing method thereof |
| KR20180038030A (en) | 2015-08-21 | 2018-04-13 | 신닛테츠스미킨 카부시키카이샤 | Steel plate |
| KR20180031738A (en) | 2015-08-31 | 2018-03-28 | 신닛테츠스미킨 카부시키카이샤 | Steel plate |
| US11519061B2 (en) | 2015-08-31 | 2022-12-06 | Nippon Steel Corporation | Steel sheet |
| WO2017051998A1 (en) * | 2015-09-22 | 2017-03-30 | 현대제철 주식회사 | Plated steel plate and manufacturing method thereof |
| US10941460B2 (en) | 2015-09-22 | 2021-03-09 | Hyundai Steel Company | Plated steel plate and manufacturing method thereof |
| US11453927B2 (en) | 2017-02-13 | 2022-09-27 | Jfe Steel Corporation | Cold rolled steel sheet and method of manufacturing the same |
| WO2018147211A1 (en) | 2017-02-13 | 2018-08-16 | Ｊｆｅスチール株式会社 | Cold rolled steel sheet and method for manufacturing same |
| KR20190107693A (en) | 2017-02-13 | 2019-09-20 | 제이에프이 스틸 가부시키가이샤 | Cold rolled steel sheet and its manufacturing method |
| WO2019093689A1 (en) * | 2017-11-07 | 2019-05-16 | 주식회사 포스코 | High-strength, low-toughness cold-rolled steel sheet having excellent fracture properties, and method for producing same |
| CN114207172A (en) * | 2019-07-31 | 2022-03-18 | 杰富意钢铁株式会社 | High-strength steel sheet, high-strength member, and method for producing same |
| CN114207172B (en) * | 2019-07-31 | 2023-07-18 | 杰富意钢铁株式会社 | High-strength steel sheet, high-strength member, and method for producing same |
| KR20220127894A (en) | 2020-01-22 | 2022-09-20 | 닛폰세이테츠 가부시키가이샤 | Steel plate and its manufacturing method |
| WO2021149676A1 (en) | 2020-01-22 | 2021-07-29 | 日本製鉄株式会社 | Steel sheet and method for producing same |
| WO2023052814A1 (en) * | 2021-09-29 | 2023-04-06 | Arcelormittal | Cold rolled and heat treated steel sheet and a method of manufacturing thereof |
Also Published As
| Publication number | Publication date |
|---|---|
| JP4995109B2 (en) | 2012-08-08 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| JP4995109B2 (en) | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same | |
| EP2757169B1 (en) | High-strength steel sheet having excellent workability and method for producing same | |
| JP5088023B2 (en) | High-strength cold-rolled steel sheet with excellent workability and method for producing the same | |
| JP6354919B1 (en) | Thin steel plate and manufacturing method thereof | |
| KR101621639B1 (en) | Steel sheet, plated steel sheet, method for producing steel sheet, and method for producing plated steel sheet | |
| JP6696209B2 (en) | High strength steel sheet manufacturing method | |
| JP5347738B2 (en) | Method for producing precipitation strengthened cold rolled steel sheet | |
| US20120175028A1 (en) | High strength steel sheet and method for manufacturing the same | |
| JP6597374B2 (en) | High strength steel plate | |
| JP5157375B2 (en) | High-strength cold-rolled steel sheet excellent in rigidity, deep drawability and hole expansibility, and method for producing the same | |
| WO2013150669A1 (en) | Galvannealed hot-rolled steel sheet and method for manufacturing same | |
| JP2009209451A (en) | High-strength hot-dip galvanized steel sheet excellent in workability and process for production thereof | |
| JP4959418B2 (en) | High-strength cold-rolled steel sheet and manufacturing method thereof | |
| WO2013084478A1 (en) | Method for manufacturing high-strength cold-rolled steel sheet having excellent aging resistance and bake hardenability | |
| JP6601253B2 (en) | High strength steel plate | |
| JP5712771B2 (en) | Steel sheet with excellent Young's modulus in the direction perpendicular to rolling and method for producing the same | |
| JP5088041B2 (en) | High Young's modulus, high strength cold-rolled steel sheet with excellent local ductility and method for producing the same | |
| WO2011118421A1 (en) | Method for producing high-strength steel plate having superior deep drawing characteristics | |
| JP5151468B2 (en) | High-strength cold-rolled steel sheet excellent in workability and impact resistance and method for producing the same | |
| JP6696208B2 (en) | High strength steel sheet manufacturing method | |
| US11643701B2 (en) | High-strength hot-dip galvanized steel sheet and manufacturing method therefor | |
| JP2019044269A (en) | High intensity cold rolled thin steel sheet | |
| JP6384623B2 (en) | High strength steel plate and manufacturing method thereof | |
| JP6417841B2 (en) | Cold rolled steel sheet manufacturing method | |
| JP5440375B2 (en) | Hot-dip galvanized steel sheet and manufacturing method thereof |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20100209 |
|
| A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20120213 |
|
| A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20120221 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20120328 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20120417 |
|
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20120509 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20150518 Year of fee payment: 3 |
|
| R151 | Written notification of patent or utility model registration |
Ref document number: 4995109 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R151 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20150518 Year of fee payment: 3 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20150518 Year of fee payment: 3 |
|
| S533 | Written request for registration of change of name |
Free format text: JAPANESE INTERMEDIATE CODE: R313533 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20150518 Year of fee payment: 3 |
|
| R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |
|
| S533 | Written request for registration of change of name |
Free format text: JAPANESE INTERMEDIATE CODE: R313533 |
|
| R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |