JP2007146275A - Low yield ratio type steel sheet with high young's modulus, hot-dip galvanized steel sheet, galvannealed steel sheet and steel tube, and their manufacturing method - Google Patents
Low yield ratio type steel sheet with high young's modulus, hot-dip galvanized steel sheet, galvannealed steel sheet and steel tube, and their manufacturing method Download PDFInfo
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本発明は、低降伏比型高ヤング率鋼板、溶融亜鉛めっき鋼板、合金化溶融亜鉛めっき鋼板及び鋼管並びにそれらの製造方法に関するものである。 The present invention relates to a low yield ratio type high Young's modulus steel plate, a hot dip galvanized steel plate, an alloyed hot dip galvanized steel plate, a steel pipe, and methods for producing them.
鋼材のヤング率を高める技術についてはこれまでにも多数の報告がある。そのほとんどが、圧延方向（ＲＤ）に対して直角な方向（ＴＤ）のヤング率を高める技術に関するものである。
以下の特許文献１〜９などは、いずれも圧延をα＋γ２相域（フェライト相とオーステナイト相が共存する温度域）で行うことによってＴＤ方向のヤング率を高める技術を開示している。
また、以下の特許文献１０は表層にＡｒ３変態点未満での圧延を加えることによってＴＤ方向のヤング率を高める技術を開示している。
一方、ＴＤ方向のヤング率と同時にＲＤ方向のヤング率を高める技術に関する開示もある。すなわち、特許文献１１は一定方向への圧延に加えてそれと直角方向の圧延を施すことで両方のヤング率を高めるものである。しかしながら、薄板の連続熱延プロセスにおいては、圧延方向を途中で変化することは生産性を著しく阻害するため、現実的ではない。
There have been many reports on techniques for increasing the Young's modulus of steel. Most of them relate to techniques for increasing the Young's modulus in the direction (TD) perpendicular to the rolling direction (RD).
The following Patent Documents 1 to 9 disclose techniques for increasing the Young's modulus in the TD direction by performing rolling in an α + γ2 phase region (a temperature region in which a ferrite phase and an austenite phase coexist).
Patent Document 10 below discloses a technique for increasing the Young's modulus in the TD direction by applying rolling below the Ar3 transformation point to the surface layer.
On the other hand, there is also a disclosure relating to a technique for increasing the Young's modulus in the RD direction simultaneously with the Young's modulus in the TD direction. In other words, Patent Document 11 increases both Young's moduli by rolling in a direction perpendicular to the rolling in addition to rolling in a certain direction. However, in the continuous hot rolling process for thin plates, changing the rolling direction in the middle significantly impairs productivity, and is not realistic.
また、特許文献１２はヤング率の高い冷延鋼板に関する技術を開示しているが、これもＴＤ方向のヤング率は高いが、ＲＤ方向のヤング率が高いわけではない。
更に、特許文献１３はＭｏ，Ｎｂ，Ｂを複合添加してヤング率を向上させる技術を開示しているが、熱延条件がまったく異なるため、ＴＤ方向のヤング率は高いが、ＲＤ方向のヤング率が高いわけではない。
また、本発明者らの一部はＲＤ方向のヤング率が高い鋼板を特許文献１４に提案した。しかし、これは鋼板の降伏比の低下を課題としたものではなく、ミクロ組織及び熱延後の冷却条件については開示していない。
Furthermore, Patent Document 13 discloses a technique for improving the Young's modulus by adding Mo, Nb, and B in combination, but because the hot rolling conditions are completely different, the Young's modulus in the TD direction is high, but the Young in the RD direction is high. The rate is not high.
Further, some of the present inventors have proposed a steel sheet having a high Young's modulus in the RD direction in Patent Document 14. However, this is not intended to reduce the yield ratio of the steel sheet, and does not disclose the microstructure and the cooling conditions after hot rolling.
上述の通り、従来にも高ヤング率鋼板と称するものは存在したが、いずれも圧延方向と直角方向（幅方向）のヤング率が高い鋼板であった。ところが鋼板の幅は最大でも２ｍ程度であり、ヤング率最大の方向を部材の長手方向とする場合には、その長さを幅以上にすることはできなかった。したがって、長物部材に対しては圧延方向のヤング率が高い鋼板が切望されていた。また、製造法についても圧延反力の変動しやすいα+γ域での熱延が前提となっており生産性に問題があった。 As described above, there has been a steel sheet having a high Young's modulus in the past, but all of them were steel sheets having a high Young's modulus in the direction perpendicular to the rolling direction (width direction). However, the width of the steel sheet is about 2 m at the maximum, and when the direction with the maximum Young's modulus is the longitudinal direction of the member, the length cannot be made larger than the width. Therefore, a steel plate having a high Young's modulus in the rolling direction has been desired for long members. In addition, the production method is premised on hot rolling in the α + γ region where the rolling reaction force is likely to fluctuate, and there is a problem in productivity.
また、特に自動車用鋼板において、乗車人員の安全性確保の及び燃費向上の観点から、高強度鋼板の使用比率が拡大している。鋼板を自動車用や建材用の部品に加工する際、形状凍結性が大きな問題となる。例えば曲げ加工を行った後、荷重が除荷される際に鋼板が元の形状に戻ろうとするスプリングバック現象が起こり、所望の形状が得られない現象である。この現象は高強度化に伴い顕在化するため、高強度鋼板を部材に適用する際の障害となっている。 In particular, in steel sheets for automobiles, the use ratio of high-strength steel sheets is increasing from the viewpoint of ensuring the safety of passengers and improving fuel consumption. When processing steel sheets into parts for automobiles and building materials, shape freezing becomes a big problem. For example, after a bending process, when the load is unloaded, a springback phenomenon occurs in which the steel sheet returns to its original shape, and a desired shape cannot be obtained. Since this phenomenon becomes apparent as the strength increases, it becomes an obstacle when a high strength steel sheet is applied to a member.
本発明者らは、上記の目標を達成するために、γ単相域圧延を前提として鋭意、研究を遂行し、以下に述べるような従来にはない知見を得た。
即ち、Ｍｏ，Ｎｂ，Ｂ，Ｔｉを所定量含有する鋼の表面近傍に所定の集合組織を発達せしめることによって圧延方向のヤング率が高い鋼板を発明することに成功したものである。また、本発明によって得られる鋼板は表面近傍では２４０ＧＰａ以上の特に高いヤング率が得られることから曲げ剛性が著しく向上し、例えば形状凍結性も著しく改善される。
In order to achieve the above-mentioned goal, the present inventors diligently conducted research on the premise of γ single-phase rolling, and obtained the following unprecedented knowledge.
That is, the inventors succeeded in inventing a steel sheet having a high Young's modulus in the rolling direction by developing a predetermined texture near the surface of the steel containing a predetermined amount of Mo, Nb, B, and Ti. In addition, since the steel sheet obtained by the present invention has a particularly high Young's modulus of 240 GPa or more near the surface, the bending rigidity is remarkably improved, for example, the shape freezing property is remarkably improved.
鋼材の高強度化に伴いスプリングバックなどの形状凍結不良の度合いが大きくなる要因は、プレス変形の際にかけられた荷重が除荷された際の戻り量が大きいことにある。したがって、ヤング率を高くすれば、戻り量を抑え、スプリングバックを低減することが可能になる。加えて曲げ変形時には曲げモーメントの大きい表層付近の変形挙動が形状凍結性に著しい影響を及ぼすことから、表層だけのヤング率を向上させることで、著しい改善が可能である。
また、降伏比（引張強さに対する降伏強さの比）を下げる事で、衝突時の吸収エネルギーが増加し、衝突安全性の向上に大きく寄与する。このためフェライト又はベイナイトを体積分率で最大の組織とし、体積分率で２〜２５％のマルテンサイト組織を含む複合組織鋼とする事で、降伏強さの低下と伴に引張延性が向上し、吸収エネルギーが増加する。
A factor that increases the degree of shape freezing failure such as springback as the strength of steel is increased is that the return amount when the load applied during press deformation is unloaded is large. Therefore, if the Young's modulus is increased, the return amount can be suppressed and the spring back can be reduced. In addition, since the deformation behavior near the surface layer having a large bending moment greatly affects the shape freezing property during bending deformation, significant improvement is possible by increasing the Young's modulus of the surface layer alone.
Also, by reducing the yield ratio (ratio of yield strength to tensile strength), the absorbed energy at the time of collision increases, greatly contributing to the improvement of collision safety. For this reason, by making ferrite or bainite the maximum structure in volume fraction and making it a composite structure steel containing a martensite structure with a volume fraction of 2 to 25%, the tensile ductility is improved with a decrease in yield strength. , The absorbed energy increases.
本発明は、このような思想と新知見に基づいて構築された従来にはない全く新しい鋼板及びその製造方法であり、その要旨とするところは以下の通りである。
（１）本発明の低降伏比型高ヤング率鋼板は、質量％で、Ｃ：０．０００５〜０．３０％、Ｓｉ：２．５％以下、Ｍｎ：０．１〜５．０％、Ｐ：０．１５％以下、Ｓ：０．０１５％以下、Ａｌ：０．１５％以下、Ｎ：０．０１％以下、及びＭｏ：０．００５〜１．５％、Ｎｂ：０．００５〜０．２０％ 、Ｔｉ：４８／１４×Ｎ［ｍａｓｓ％］以上，０．２％以下、Ｂ：０．０００１〜０．０１％のいずれか１種又は２種以上を合計で０．０１５〜１．９１％含有し、残部鉄及び不可避的不純物からなり、フェライト又はベイナイトを体積分率最大の組織とし、体積分率で２〜２５％のマルテンサイトを含む複合組織鋼であり、かつ板厚の１／８層における｛１１０｝＜２２３＞、｛１１０｝＜１１１＞の一方又は双方の極密度が１０以上であり、圧延方向のヤング率が２３０ＧＰａ超であることを特徴とする。
（２）本発明の低降伏比型高ヤング率鋼板は、（１）において選択的に含有する組成を全て含み、Ｍｏ：０．１〜１．５％、Ｂ：０．０００６〜０．０１％、Ｎｂ：０．０１〜０．２０％、Ｔｉ：４８／１４×Ｎ［ｍａｓｓ％］以上，０．２％以下を含有し、板厚の１／８層における｛１１０｝＜００１＞の極密度が６以下であることを特徴とする。
（３）本発明の低降伏比型高ヤング率鋼板は、前記（１）又は（２）に記載の低降伏比型高ヤング率鋼板において、Ｃａ：０．０００５〜０．０１質量％を含むことを特徴とする。
（４）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（３）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、Ｓｎ，Ｃｏ，Ｚｎ，Ｗ，Ｚｒ，Ｖ，Ｍｇ，Ｒｅｍの１種又は２種以上を合計で０．００１〜１．０質量％含むことを特徴とする。
（５）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（４）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、Ｎｉ，Ｃｕ，Ｃｒの１種又は２種以上を合計で０．００１〜４．０質量％含むことを特徴とする。
The present invention is a completely new steel sheet and a method for manufacturing the same that have been constructed based on these ideas and new findings. The gist of the present invention is as follows.
(1) The low yield ratio type high Young's modulus steel sheet of the present invention is in mass%, C: 0.0005 to 0.30%, Si: 2.5% or less, Mn: 0.1 to 5.0%, P: 0.15% or less, S: 0.015% or less, Al: 0.15% or less, N: 0.01% or less, and Mo: 0.005 to 1.5%, Nb: 0.005 0.20%, Ti: 48/14 × N [mass%] or more, 0.2% or less, B: 0.0001 to 0.01% of any one or two or more in total 0.015 1.91% content, consisting of the balance iron and inevitable impurities, ferrite or bainite with a maximum volume fraction structure, a composite structure steel containing martensite with a volume fraction of 2 to 25%, and a plate thickness The pole density of one or both of {110} <223> and {110} <111> in the 1/8 layer of , The rolling direction of the Young's modulus is equal to or is 230GPa greater.
(2) The low yield ratio type high Young's modulus steel sheet of the present invention includes all the compositions selectively contained in (1), Mo: 0.1 to 1.5%, and B: 0.0006 to 0.01. %, Nb: 0.01 to 0.20%, Ti: 48/14 × N [mass%] or more and 0.2% or less, and {110} <001> in 1/8 layer of the plate thickness The pole density is 6 or less.
(3) The low yield ratio type high Young's modulus steel sheet of the present invention includes Ca: 0.0005 to 0.01% by mass in the low yield ratio type high Young's modulus steel sheet according to (1) or (2). It is characterized by that.
(4) The low yield ratio type high Young's modulus steel sheet of the present invention is the low yield ratio type high Young's modulus steel sheet according to any one of (1) to (3), wherein Sn, Co, Zn, W, One or more of Zr, V, Mg, and Rem are included in a total amount of 0.001 to 1.0 mass%.
(5) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (4) above, which is one of Ni, Cu, and Cr. Or 2 or more types are included in total 0.001-4.0 mass%, It is characterized by the above-mentioned.
（６）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（５）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、板厚の１／８層における｛１１０｝＜００１＞の極密度が３以下であることを特徴とする。
（７）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（６）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、少なくとも板厚の表層から１／８層における圧延方向のヤング率が２４０ＧＰａ以上であることを特徴とする。
（８）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（７）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、更に、板厚１／２層における｛２１１｝＜０１１＞の極密度が６以上であることを特徴とする。
（９）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（８）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、更に、板厚１／２層における｛３３２｝＜１１３＞の極密度が６以上であることを特徴とする。
（１０）本発明の低降伏比型高ヤング率鋼板は、前記（１）〜（９）のいずれか１項に記載の低降伏比型高ヤング率鋼板において、更に、板厚１／２層における｛１００｝＜０１１＞の極密度が６以下であることを特徴とする。
(6) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (5), wherein the thickness is 1/8 layer. The pole density of {110} <001> is 3 or less.
(7) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of the above (1) to (6). The Young's modulus in the rolling direction in the eight layers is 240 GPa or more.
(8) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (7), further having a thickness of 1/2 layer. {211} <011> has a pole density of 6 or more.
(9) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (8), further comprising a thickness 1/2 layer. The pole density of {332} <113> in is 6 or more.
(10) The low yield ratio type high Young's modulus steel plate of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (9), further having a thickness of 1/2 layer. {100} <011> has a pole density of 6 or less.
（１１）本発明の低降伏比型高ヤング率溶融亜鉛メッキ鋼板は、前記（１）〜（１０）の何れか１項に記載の低降伏比型高ヤング率鋼板に、溶融亜鉛めっきが施されていることを特徴とする。
（１２）本発明の低降伏比型高ヤング率合金化溶融亜鉛メッキ鋼板は、前記（１）〜（１０）の何れか１項に記載の低降伏比型高ヤング率鋼板に、合金化溶融亜鉛めっきが施されていることを特徴とする。
（１３）本発明の低降伏比型高ヤング率鋼管は、前記（１）〜（１０）の何れか１項に記載の低降伏比型高ヤング率鋼板、前記（１１）記載の低降伏比型高ヤング率溶融亜鉛めっき鋼板、又は前記（１２）記載の低降伏比型高ヤング率合金化溶融亜鉛めっき鋼板が、任意の方向に巻かれていることを特徴とする。
（１４）本発明の低降伏比型高ヤング率鋼板の製造方法は、前記（１）〜（１０）の何れか１項に記載の高ヤング率鋼板を製造する方法であって、前記（１）〜（５）のいずれか１項に記載の化学成分を有するスラブを１０００℃以上の温度に加熱し、熱間圧延をする際、圧延ロールと鋼板との摩擦係数が０．２超、以下の［１］式で計算される有効ひずみ量ε＊が０．４以上、かつ圧下率の合計が５０％以上となるように圧延を行い、Ａｒ３変態点［℃］以上９００℃以下の温度で熱間圧延を終了し、３０ｓ以内の空冷を行った後、５〜１５０℃／ｓの冷却速度で２５〜３００℃まで冷却し、巻き取ることを特徴とする。
(11) The low yield ratio high Young's modulus hot dip galvanized steel sheet of the present invention is obtained by subjecting the low yield ratio high Young's modulus steel sheet according to any one of (1) to (10) to hot dip galvanization. It is characterized by being.
(12) The low yield ratio type high Young's modulus alloyed hot-dip galvanized steel sheet of the present invention is alloyed and melted to the low yield ratio type high Young's modulus steel sheet according to any one of (1) to (10). It is characterized by being galvanized.
(13) The low yield ratio type high Young's modulus steel pipe of the present invention is the low yield ratio type high Young's modulus steel plate according to any one of (1) to (10), and the low yield ratio according to (11). The type high Young's modulus hot-dip galvanized steel sheet or the low yield ratio type high Young's modulus alloyed hot-dip galvanized steel sheet described in (12) is wound in any direction.
(14) The method for producing a low yield ratio type high Young's modulus steel sheet according to the present invention is a method for producing the high Young's modulus steel sheet according to any one of (1) to (10), wherein the (1 ) When the slab having the chemical component according to any one of (5) is heated to a temperature of 1000 ° C. or higher and hot-rolled, the friction coefficient between the rolling roll and the steel sheet is more than 0.2, below The rolling is performed so that the effective strain amount ε * calculated by the formula [1] is 0.4 or more and the total reduction ratio is 50% or more, and at a temperature of Ar3 transformation point [° C] or more and 900 ° C or less. After the hot rolling is finished and air cooling is performed within 30 s, it is cooled to 25 to 300 ° C. at a cooling rate of 5 to 150 ° C./s and wound up.
上記の式１において、ｎは仕上げ熱延の圧延スタンド数、εｊはｊ番目のスタンドで加えられたひずみ、εｎはｎ番目のスタンドで加えられたひずみ、ｔｉはｉ〜ｉ＋１番目のスタンド間の走行時間［ｓ］、τｉは気体常数Ｒ（＝１．９８７）とｉ番目のスタンドの圧延温度Ｔｉ［Ｋ］によって以下の［２］式で計算できる。 In Equation 1 above, n is the number of rolling stands for finish hot rolling, εj is the strain applied at the jth stand, εn is the strain applied at the nth stand, and ti is between i to i + 1th stands. The traveling time [s] and τi can be calculated by the following equation [2] according to the gas constant R (= 1.987) and the rolling temperature Ti [K] of the i-th stand.
（１５）本発明の低降伏比型高ヤング率鋼板の製造方法は、前記（１）〜（１０）の何れか１項に記載の高ヤング率鋼板を製造する方法であって、前記（１）〜（５）のいずれか１項に記載の化学成分を有するスラブを１０００℃以上の温度に加熱し、熱間圧延をする際、圧延ロールと鋼板との摩擦係数が０．２超、以下の［３］式で計算される有効ひずみ量ε＊が０．４以上、かつ圧下率の合計が５０％以上となるように圧延を行い、Ａｒ３変態点［℃］以上９００℃以下の温度で熱間圧延を終了し、酸洗後に１０％以上６０％未満の圧下率で冷間圧延を施し、最高加熱温度Ａｃ１変態点［℃］以上０．５×（Ａｃ１＋Ａｃ３）［℃］以下の温度範囲で焼鈍後、１〜１５０℃／ｓの冷却速度にて２５〜３８０℃まで冷却することを特徴とする。 (15) The method for producing a low yield ratio type high Young's modulus steel sheet according to the present invention is a method for producing the high Young's modulus steel sheet according to any one of (1) to (10), wherein the (1 ) When the slab having the chemical component according to any one of (5) is heated to a temperature of 1000 ° C. or higher and hot-rolled, the friction coefficient between the rolling roll and the steel sheet is more than 0.2, below The rolling is performed so that the effective strain amount ε * calculated by the formula [3] is 0.4 or more and the total rolling reduction is 50% or more, and at a temperature of Ar3 transformation point [° C] or more and 900 ° C or less. After the hot rolling is finished, the steel sheet is cold-rolled at a reduction rate of 10% or more and less than 60% after pickling, and the maximum heating temperature Ac1 transformation point [° C.] to 0.5 × (Ac1 + Ac3) [° C.] temperature range After annealing, it is cooled to 25 to 380 ° C. at a cooling rate of 1 to 150 ° C./s. .
上記の式３において、ｎは仕上げ熱延の圧延スタンド数、εｊはｊ番目のスタンドで加えられたひずみ、εｎはｎ番目のスタンドで加えられたひずみ、ｔｉはｉ〜ｉ＋１番目のスタンド間の走行時間［ｓ］、τｉは気体常数Ｒ（＝１．９８７）とｉ番目のスタンドの圧延温度Ｔｉ［Ｋ］によって以下の［４］式で計算できる。 In the above formula 3, n is the number of rolling hot rolling stands, εj is the strain applied at the jth stand, εn is the strain applied at the nth stand, and ti is between i and i + 1th stands. The traveling time [s] and τi can be calculated by the following equation [4] based on the gas constant R (= 1.987) and the rolling temperature Ti [K] of the i-th stand.
（１６）本発明の低降伏比型高ヤング率鋼板の製造方法は、前記（１４）又は（１５）記載の低降伏比型高ヤング率鋼板の製造方法において、熱間圧延を実施する際に異周速率が１％以上の異周速圧延を少なくとも１パス以上施すことを特徴とする。
（１７）本発明の低降伏比型高ヤング率鋼板の製造方法は、前記（１４）〜（１６）の何れか１項に記載の低降伏比型高ヤング率鋼板の製造方法において、熱間圧延を実施する際にロール径が７００ｍｍ以下の圧延ロールを少なくとも１つ以上使用することを特徴とする。
（１８）本発明の低降伏比型高ヤング率鋼板の製造方法は、前記（１４）、（１６）、（１７）の何れか１項に記載の方法により製造した熱延鋼板を酸洗後、１０％以上６０％未満の圧下率で冷間圧延を施した後に、最高加熱温度が５００℃以上０．５×（Ａｃ１＋Ａｃ３）［℃］以下の温度範囲で焼鈍することを特徴とする。
(16) The method for producing a low yield ratio type high Young's modulus steel sheet of the present invention is the method for producing a low yield ratio type high Young's modulus steel sheet described in (14) or (15) above, when hot rolling is performed. It is characterized in that at least one pass of different circumferential speed rolling with a different circumferential speed ratio of 1% or more is performed.
(17) The method for producing a low yield ratio type high Young's modulus steel sheet according to the present invention is a method for producing a low yield ratio type high Young's modulus steel sheet according to any one of the above (14) to (16). When rolling, at least one rolling roll having a roll diameter of 700 mm or less is used.
(18) The method for producing a low yield ratio type high Young's modulus steel sheet according to the present invention is after pickling a hot-rolled steel sheet produced by the method according to any one of (14), (16), and (17). After cold rolling at a rolling reduction of 10% or more and less than 60%, annealing is performed in a temperature range where the maximum heating temperature is 500 ° C. or more and 0.5 × (Ac1 + Ac3) [° C.] or less.
（１９）本発明の低降伏比型高ヤング率溶融亜鉛めっき鋼板の製造方法は、前記（１１）記載の低降伏比型高ヤング率溶融亜鉛めっき鋼板を製造する方法であって、前記（１４）〜（１８）の何れか１項に記載の方法で製造した低降伏比型高ヤング率鋼板に、溶融亜鉛めっきを施すことを特徴とする。 (19) The method for producing a low yield ratio type high Young's modulus hot-dip galvanized steel sheet according to the present invention is a method for producing the low yield ratio type high Young's modulus hot-dip galvanized steel sheet described in (11) above, The low yield ratio type high Young's modulus steel plate produced by the method according to any one of (1) to (18) is subjected to hot dip galvanization.
（２０）本発明の低降伏比型高ヤング率溶融亜鉛めっき鋼板の製造方法は、前記（１１）記載の低降伏比型高ヤング率溶融亜鉛めっき鋼板を製造する方法であって、前記（１６）〜（１９）のいずれか１項に記載の低降伏比高ヤング率鋼板の製造方法に引き続き、更に、連続ラインにて、溶融亜鉛めっきを施すことを特徴とする。 (20) The method for producing a low yield ratio type high Young's modulus hot-dip galvanized steel sheet according to the present invention is a method for producing the low yield ratio type high Young's modulus hot-dip galvanized steel sheet described in (11) above. ) To (19), the method for producing a steel plate with a low yield ratio and a high Young's modulus is characterized in that hot dip galvanizing is further performed in a continuous line.
（２１）本発明の低降伏比型高ヤング率合金化溶融亜鉛めっき鋼板の製造方法は、前記（１２）記載の低降伏比型高ヤング率合金化溶融亜鉛めっき鋼板を製造する方法であって、前記（１９）又は（２０）記載の溶融亜鉛めっきを施した後、４５０〜６００℃までの温度範囲で５ｓ以上の熱処理を行うことを特徴とする。
（２２）本発明の低降伏比型高ヤング率鋼管の製造方法は、前記（１３）記載の鋼管を製造する方法であって、前記（１４）〜（１８）のいずれか１項に記載の製造方法により得られた低降伏比型高ヤング率鋼板、前記（１９）又は（２０）記載の製造方法により得られた低降伏比型高ヤング率溶融亜鉛めっき鋼板又は前記（２１）記載の製造方法により得られた低降伏比型高ヤング率合金化溶融亜鉛めっき鋼板を任意の方向に巻いて鋼管にすることを特徴とする。
(21) A method for producing a low yield ratio high Young's modulus galvannealed steel sheet according to the present invention is a method for producing a low yield ratio high Young's modulus galvannealed steel sheet according to (12) above. After the hot dip galvanizing described in (19) or (20), a heat treatment for 5 s or more is performed in a temperature range of 450 to 600 ° C.
(22) A method for producing a low yield ratio type high Young's modulus steel pipe according to the present invention is a method for producing the steel pipe described in (13) above, and is described in any one of (14) to (18). Low yield ratio type high Young's modulus steel sheet obtained by the production method, low yield ratio type high Young's modulus hot dip galvanized steel sheet obtained by the production method described in (19) or (20) above, or production according to (21) above A low yield ratio type high Young's modulus galvannealed steel sheet obtained by the method is rolled into an arbitrary direction to form a steel pipe.
本発明によれば、特に圧延方向のヤング率に優れた低降伏比型鋼板、溶融亜鉛めっき鋼板、合金化溶融亜鉛めっき鋼板及び鋼管並びにそれらの製造方法を得ることができる。 According to the present invention, it is possible to obtain a low yield ratio steel sheet, a hot dip galvanized steel sheet, an alloyed hot dip galvanized steel sheet, a steel pipe, and their production methods that are particularly excellent in the Young's modulus in the rolling direction.
以下に本発明について最良の形態に基づいて詳細に説明するが、本発明が以下に記載する種々の実施の形態に制限されるものではないのは勿論である。
ここに、本発明において鋼組成及び製造条件を上述のように限定する理由についてさらに説明する。
まず、成分について説明する。
Ｃは安価に引張強度を増加させる元素であるので、その添加量は狙いとする強度レベルに応じて変化するが、Ｃの下限を０．０００５％とするのは製鋼技術上困難でコストアップとなるだけでなく、溶接部の疲労特性が劣化するのでこれを下限とする。一方、Ｃ量が０．３０％を超えると成形性の劣化を招いたり、溶接性を損なったりするのでこれを上限とする。
Ｓｉは固溶体強化元素として強度を増加させる働きがあることの他、マルテンサイトやベイナイトさらには残留γ等を含む組織を得るためにも有効であり、その添加量は狙いとする強度レベルに応じて変化するが、添加量が２．５％超となるとプレス成形性が劣悪となったり、化成処理性の低下を招いたりするのでこれを上限とする。
溶融亜鉛めっきを施す場合には、めっき密着性の低下、合金化反応の遅延による生産性の低下などの問題が生ずるのでＳｉは１．２％以下とすることが好ましい。下限は特に設けないが、Ｓｉを０．００１％以下とするのは製造コストが高くなるのでこれが実質的な下限である。
The present invention will be described in detail below based on the best mode, but the present invention is of course not limited to the various embodiments described below.
Here, the reason why the steel composition and production conditions are limited as described above in the present invention will be further described.
First, components will be described.
Since C is an element that increases the tensile strength at a low cost, the amount of addition varies depending on the target strength level. However, it is difficult to make the lower limit of C 0.0005% because of the steelmaking technology, which increases costs. In addition, since the fatigue characteristics of the welded portion deteriorate, this is set as the lower limit. On the other hand, if the C content exceeds 0.30%, the formability is deteriorated or the weldability is impaired.
Si has a function of increasing strength as a solid solution strengthening element, and is also effective for obtaining a structure containing martensite, bainite and residual γ, and the amount of addition depends on the target strength level. However, if the addition amount exceeds 2.5%, the press formability is deteriorated or the chemical conversion property is lowered, so this is the upper limit.
When hot dip galvanization is performed, Si is preferably 1.2% or less because problems such as a decrease in plating adhesion and a decrease in productivity due to a delay in alloying reaction occur. Although there is no particular lower limit, Si is made 0.001% or less because the manufacturing cost increases, and this is a practical lower limit.
Ｍｎはγ相を安定化し、γ域を低温まで拡張するのでγ域低温圧延を容易にする。また、表層近傍の剪断集合組織形成にＭｎ自体が有利に作用している可能性もある。これらの観点から、Ｍｎは０．１％以上添加する。この観点から０．５％以上添加することが望ましい。更に望ましくは１．５％以上添加する。一方、５．０％を超えると強度が高くなりすぎて延性が低下したり、亜鉛めっきの密着性が阻害されたりするのでこれを上限とする。好ましくは２．９〜４．０％とする。
ＰはＳｉと同様に安価に強度を高める元素として知られており強度を増加する必要がある場合にはさらに積極的に添加する。また、Ｐは熱延組織を微細にし、加工性を向上する効果も有する。ただし、添加量が０．１５％を超えると、スポット溶接後の疲労強度が劣悪となったり、降伏強度が増加し過ぎたりしてプレス時に面形状不良を引き起こす。さらに、連続溶融亜鉛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。また、２次加工性も劣化する。したがって、その上限値を０．１５％とする。
Ｓは、０．０１５％超では熱間割れの原因となったり加工性を劣化させたりするので、これを上限とする。
Since Mn stabilizes the γ phase and expands the γ region to a low temperature, it facilitates γ region low temperature rolling. In addition, Mn itself may have an advantageous effect on the formation of a shear texture near the surface layer. From these viewpoints, Mn is added by 0.1% or more. From this viewpoint, it is desirable to add 0.5% or more. More desirably, 1.5% or more is added. On the other hand, if it exceeds 5.0%, the strength becomes too high and the ductility is lowered or the adhesion of galvanization is hindered. Preferably it is set to 2.9 to 4.0%.
Like Si, P is known as an element that increases the strength at a low cost, and when it is necessary to increase the strength, P is more actively added. P also has the effect of making the hot-rolled structure fine and improving workability. However, if the addition amount exceeds 0.15%, the fatigue strength after spot welding becomes poor, or the yield strength increases excessively, causing surface shape defects during pressing. Furthermore, the alloying reaction becomes extremely slow during continuous hot dip galvanizing, and productivity is lowered. Also, the secondary workability is deteriorated. Therefore, the upper limit is set to 0.15%.
If S exceeds 0.015%, it causes hot cracking and deteriorates workability, so this is the upper limit.
Ｍｏ，Ｎｂ，Ｔｉ及びＢは本発明において重要な添加元素である。これらの元素の添加によって初めて圧延方向のヤング率を高めることが可能となる。この理由は必ずしも明らかではないが、熱延中の再結晶が抑制され、γ相の加工集合組織が鮮鋭化することで、結果的に鋼板と熱延ロールとの摩擦に起因する剪断変形集合組織にも変化が生じ、熱延板の板厚表層から板厚１／４層近傍までの範囲において、非常に先鋭な集合組織が形成され、圧延方向のヤング率が高くなる。Ｍｏ，Ｎｂ，Ｔｉ及びＢ量の下限はそれぞれＭｏ：０．００５％，Ｎｂ：０．００５％，Ｔｉ：４８／１４×Ｎ［ｍａｓｓ％］，Ｂ：０．０００１％，望ましくはＭｏ：０．０３％、Ｎｂ：０．０１％、Ｔｉ：０．０３％、Ｂ：０．０００３％、更に望ましくはＭｏ：０．１％、Ｎｂ：０．０３％、Ｔｉ：０．０５％、Ｂ：０．０００６％である。前記の下限より少ない量の添加では、上述のヤング率向上効果が小さくなってしまうからである。
一方、Ｍｏ，Ｎｂ，Ｔｉ，ＢをそれぞれＭｏ：１．５％超、Ｎｂ：０．２％超、Ｔｉ：０．２％超、Ｂ：０．０１％超添加してもヤング率の向上効果は飽和し、コストアップとなるのでこれを上限とする。また、これら元素の合計添加量が０．０１５％未満では十分なヤング率向上効果が得られないことから０．０１５％を合計添加量の下限とする。この観点から望ましくは合計で０．０３５％以上、更に望ましくは合計で０．０５％以上添加する。合計添加量の上限はそれぞれの添加量の上限の和である１．９１％とする。
Mo, Nb, Ti and B are important additive elements in the present invention. Only when these elements are added can the Young's modulus in the rolling direction be increased. The reason for this is not necessarily clear, but recrystallization during hot rolling is suppressed, and the processing texture of the γ phase is sharpened, resulting in a shear deformation texture resulting from friction between the steel sheet and the hot rolling roll. In the range from the thickness surface layer of the hot-rolled sheet to the vicinity of the 1/4 thickness, a very sharp texture is formed and the Young's modulus in the rolling direction is increased. The lower limits of the amounts of Mo, Nb, Ti and B are respectively Mo: 0.005%, Nb: 0.005%, Ti: 48/14 × N [mass%], B: 0.0001%, preferably Mo: 0 0.03%, Nb: 0.01%, Ti: 0.03%, B: 0.0003%, more preferably Mo: 0.1%, Nb: 0.03%, Ti: 0.05%, B : 0.0006%. This is because if the amount is less than the lower limit, the effect of improving the Young's modulus is reduced.
On the other hand, even if Mo, Nb, Ti, and B are added to each of Mo: more than 1.5%, Nb: more than 0.2%, Ti: more than 0.2%, and B: more than 0.01%, the Young's modulus is improved. Since the effect is saturated and the cost increases, this is the upper limit. Further, if the total addition amount of these elements is less than 0.015%, a sufficient Young's modulus improvement effect cannot be obtained, so 0.015% is made the lower limit of the total addition amount. From this point of view, the total amount is preferably 0.035% or more, and more preferably 0.05% or more in total. The upper limit of the total addition amount is 1.91% which is the sum of the upper limits of the respective addition amounts.
なお、Ｍｏ，Ｎｂ，ＴｉとＢの間には相互作用があり複合添加することで更に集合組織が強くなり、ヤング率が上昇することから、少なくとも２種以上を複合添加することがより望ましい。特にＴｉはγ高温域でＮと窒化物を形成し、ＢＮの生成を抑制するため、Ｂを添加する場合はＴｉも４８／１４×Ｎ％以上添加することが望ましい。
また、それぞれの元素がいずれもＭｏ：０．１５％以上，Ｎｂ：０．０１％以上，Ｔｉ：４８／１４×Ｎ％以上，Ｂ：０．０００６％以上添加されている場合には、集合組織が先鋭化し、特にヤング率を低減させる表層の｛１１０｝＜００１＞が減少し効果的なヤング率上昇がなされるため、高いＬ方向（圧延方向）ヤング率が達成されるので好ましい。
なお、これらの元素の同時添加によるヤング率向上効果は、Ｃとの組み合わせによってさらに助長される。したがってＣ量は、０．０１５％以上とすることが好ましい。
In addition, since there is an interaction between Mo, Nb, Ti, and B, and a composite addition further strengthens the texture and increases the Young's modulus, it is more desirable to add at least two or more. In particular, Ti forms nitrides with N in the γ high temperature region and suppresses the generation of BN. Therefore, when adding B, it is preferable to add Ti at 48/14 × N% or more.
In addition, when each element is added Mo: 0.15% or more, Nb: 0.01% or more, Ti: 48/14 × N% or more, B: 0.0006% or more, Since the structure is sharpened and the {110} <001> of the surface layer that particularly reduces the Young's modulus is reduced and the Young's modulus is effectively increased, a high L direction (rolling direction) Young's modulus is achieved, which is preferable.
In addition, the Young's modulus improvement effect by simultaneous addition of these elements is further promoted by the combination with C. Therefore, the C content is preferably 0.015% or more.
Ａｌは脱酸調整剤として使用しても良い。ただしＡｌは変態点を著しく高めるので低温γ域での圧延が困難となるので、上限を０．１５％とする。Ａｌの下限は特に限定しないが、脱酸の観点からは０．０１％以上とすることが好ましい。
ＮはＢと窒化物を形成し、Ｂの再結晶抑制効果を低減させることから０．０１％以下に抑える。この観点から望ましくは０．００５％、更に望ましくは０．００２％以下とする。Ｎの下限は特に設定しないが０．０００５％未満とすることは、コストがかかるばかりでそれほどの効果が得られないことから０．０００５％以上とすることが望ましい。
Ｃａは、脱酸元素として有用であるほか、硫化物の形態制御にも効果を奏するので、０．０００５〜０．０１％の範囲で添加しても良い。０．０００５％未満では効果が十分でなく、０．０１％超添加すると加工性が劣化するのでこの範囲とする。
Al may be used as a deoxidizer. However, since Al significantly increases the transformation point, rolling in the low temperature γ region becomes difficult, so the upper limit is made 0.15%. The lower limit of Al is not particularly limited, but is preferably 0.01% or more from the viewpoint of deoxidation.
N forms nitrides with B and reduces the recrystallization suppressing effect of B, so it is suppressed to 0.01% or less. From this viewpoint, it is preferably 0.005%, and more preferably 0.002% or less. Although the lower limit of N is not particularly set, setting it to less than 0.0005% is preferable to be 0.0005% or more because it is costly and not so effective.
Ca is useful as a deoxidizing element and also has an effect on controlling the form of sulfide, so Ca may be added in the range of 0.0005 to 0.01%. If it is less than 0.0005%, the effect is not sufficient, and if it exceeds 0.01%, the workability deteriorates, so this range is set.
これらを主成分とする鋼にＳｎ，Ｃｏ，Ｚｎ，Ｗ，Ｚｒ，Ｖ，Ｍｇ，Ｒｅｍの１種又は２種以上を合計で０．００１〜１．０％以下含有しても構わない。特にＷとＶはγ域の再結晶を抑制する効果があることからそれぞれ０．０１％以上添加することが好ましい。
Ｎｉ，Ｃｕ，Ｃｒは低温γ域圧延を行うためには有利な元素であるので、これらの１種又は２種以上を合計で０．００１〜４．０％の範囲で添加しても良い。０．００１％未満では顕著な効果が得られず、４．０％超添加すると加工性が劣化する。
The steel containing these as a main component may contain 0.001 to 1.0% or less in total of one or more of Sn, Co, Zn, W, Zr, V, Mg, and Rem. In particular, W and V are each preferably added in an amount of 0.01% or more because they have the effect of suppressing recrystallization in the γ region.
Since Ni, Cu, and Cr are elements that are advantageous for performing low-temperature γ region rolling, one or more of these may be added in a range of 0.001 to 4.0% in total. If it is less than 0.001%, a remarkable effect cannot be obtained, and if it exceeds 4.0%, workability deteriorates.
次に、本発明の低降伏比型高ヤング率鋼板の変態組織について説明する。
本発明の鋼板の組織は、フェライト、ベイナイトの一方又は双方と、マルテンサイトからなり、フェライト又はベイナイトを体積分率最大の組織とし、体積分率で２〜２５％のマルテンサイトを含む複合組織鋼とする。硬質相であるマルテンサイトを２〜２５％含み、フェライトやベイナイトといった比較的軟質な組織との複合組織鋼とする事で、引張強さに対する降伏強さの比（降伏比）を低くする事ができる。このとき、フェライト又はベイナイトを体積率で最大としたのは、これらの組織分率が小さいと十分な低降伏比が得られない事と、引張延性が十分得られず成形性を阻害する原因となる事から、これらを最大分率の組織とした。またマルテンサイト分率が２％未満もしくは２５％超では、十分な低降伏比が得らないため上記範囲を設けた。このときマルテンサイト分率は５〜２０％程度が好ましい。
本発明によって得られる鋼板の組織は、フェライト又はベイナイトを主相（体積分率で最大の相）とするが、両相が混在していても構わない。また、マルテンサイト組織の体積分率は２〜２５％と規定するが、オーステナイト、炭化物、窒化物を初めとする化合物が１０％以下、好ましくは３％未満存在していても構わない。すなわち要求特性に応じて組織を作り分ければ良い。
次に本発明の低降伏比型高ヤング率鋼板の集合組織、ヤング率について説明する。
板厚１／８層における｛１１０｝＜２２３＞、｛１１０｝＜１１１＞方位の一方又は双方の極密度は１０以上とする。これによって圧延方向のヤング率を高めることが可能となり、逆に１０未満では圧延方向のヤング率を２３０ＧＰａ超とすることは困難である。好ましくは１４以上、さらに好ましくは２０以上である。
これらの方位の極密度（Ｘ線ランダム強度比）は、Ｘ線回折によって測定される｛１１０｝，｛１００｝，｛２１１｝，｛３１０｝極点図のうち複数の極点図を基に級数展開法で計算した３次元集合組織（ＯＤＦ）から求めればよい。すなわち、各結晶方位の極密度を求めるには、３次元集合組織のφ２＝４５°断面における（１１０）［２−２３］、（１１０）［１−１１］の強度で代表させる。
Next, the transformation structure of the low yield ratio type high Young's modulus steel sheet of the present invention will be described.
The structure of the steel sheet of the present invention is composed of one or both of ferrite and bainite and martensite, and has a structure with the largest volume fraction of ferrite or bainite, and a composite structure steel containing martensite with a volume fraction of 2 to 25%. And The ratio of the yield strength to the tensile strength (yield ratio) can be lowered by using a composite structure steel containing 2-25% martensite, which is a hard phase, and a relatively soft structure such as ferrite and bainite. it can. At this time, the reason why the ferrite or bainite is maximized in volume ratio is that when these structural fractions are small, a sufficiently low yield ratio cannot be obtained, and the tensile ductility cannot be obtained sufficiently and the formability is hindered. Therefore, these were the organizations with the highest fractions. Further, when the martensite fraction is less than 2% or more than 25%, a sufficiently low yield ratio cannot be obtained, so the above range is provided. At this time, the martensite fraction is preferably about 5 to 20%.
The structure of the steel sheet obtained by the present invention has ferrite or bainite as the main phase (the maximum volume fraction), but both phases may be mixed. Further, the volume fraction of the martensite structure is defined as 2 to 25%, but compounds such as austenite, carbide, and nitride may be present at 10% or less, preferably less than 3%. That is, it is only necessary to create an organization according to required characteristics.
Next, the texture and Young's modulus of the low yield ratio type high Young's modulus steel sheet of the present invention will be described.
The pole density of one or both of the {110} <223> and {110} <111> orientations in the plate thickness 1/8 layer is 10 or more. This makes it possible to increase the Young's modulus in the rolling direction. Conversely, if it is less than 10, it is difficult to make the Young's modulus in the rolling direction over 230 GPa. Preferably it is 14 or more, More preferably, it is 20 or more.
The pole density (X-ray random intensity ratio) of these orientations is developed in series based on a plurality of pole figures among {110}, {100}, {211}, {310} pole figures measured by X-ray diffraction. What is necessary is just to obtain | require from the three-dimensional texture (ODF) calculated by the method. That is, in order to obtain the pole density of each crystal orientation, it is represented by the intensity of (110) [2-23] and (110) [1-11] in the φ2 = 45 ° cross section of the three-dimensional texture.
上記の極密度に関する限定は少なくとも板厚１／８層については満足し、実際には１／８層のみならず、板厚表層から１／４層までの広い範囲で成り立つことが好ましい。
さらに板厚１／８層における｛１１０｝＜００１＞（上記ＯＤＦのφ２＝４５°断面における（１１０）［００１］）方位の極密度は６以下とすることが好ましい。この方位は圧延方向のヤング率を著しく低下させることから、この方位の極密度が６超になると圧延方向のヤング率が２３０ＧＰａを超えることが困難になる。この観点から好ましくは３以下、さらに好ましくは１．５未満である。
板厚１／２層における｛２１１｝＜０１１＞（上記ＯＤＦのφ２＝４５°断面における（１１２）［１−１０］）の極密度は６以上であることが好ましい。この方位が発達すると圧延方向に対して直角の方向（以下、ＴＤ方向とする）に＜１１１＞方位が集積するためＴＤ方向のヤング率が高くなる。この極密度が６未満ではＴＤ方向のヤング率を２３０ＧＰａ超とするのは困難であるので、これを下限とする。好ましくは極密度が８以上、さらに好ましくは１０以上とする。
また、板厚１／２層における｛３３２｝＜１１３＞（上記ＯＤＦのφ２＝４５°断面における（３３２）［−１−１３］）の極密度は圧延方向のヤング率には若干の寄与が期待できる。したがってこの方位の極密度は６以上であることが好ましい。この観点から好ましくは極密度が８以上、更に好ましくは１０以上とする。
更に、板厚１／２層における｛１００｝＜０１１＞（上記ＯＤＦのφ２＝４５°断面における（００１）［１−１０］）の極密度は４５°方向のヤング率を著しく低下させることから極密度を６以下にすることが望ましい。この観点から好ましくは極密度を３以下、更に望ましくは極密度を１．５以下とする。
なお、以上で述べた結晶方位はいずれも±２．５°以内のばらつきは許容するものである。
上記の結晶方位の極密度の条件は、少なくとも１つを満足することが好ましく、複数を満足することがさらに好ましく、最適条件は、全てを満足することである。
The above-mentioned limitation on the pole density is satisfied at least for the 1 / 8th layer, and in practice, it is preferably established in a wide range from the 1 / 8th layer to the 1 / 4th layer.
Furthermore, the pole density in the {110} <001> ((110) [001] in the φ2 = 45 ° cross section of the ODF) in the 1/8 layer thickness is preferably 6 or less. Since this orientation significantly reduces the Young's modulus in the rolling direction, it is difficult for the Young's modulus in the rolling direction to exceed 230 GPa when the pole density in this orientation exceeds 6. From this viewpoint, it is preferably 3 or less, more preferably less than 1.5.
It is preferable that the pole density of {211} <011> ((112) [1-10] in the φ2 = 45 ° cross section of the ODF) in the plate thickness ½ layer is 6 or more. When this orientation develops, the <111> orientation accumulates in a direction perpendicular to the rolling direction (hereinafter referred to as the TD direction), so the Young's modulus in the TD direction increases. If the pole density is less than 6, it is difficult to make the Young's modulus in the TD direction exceed 230 GPa, so this is the lower limit. Preferably, the pole density is 8 or more, more preferably 10 or more.
Moreover, the pole density of {332} <113> in the 1/2 layer thickness ((332) [-1-13] in the φ2 = 45 ° cross section of the ODF) contributes slightly to the Young's modulus in the rolling direction. I can expect. Accordingly, the pole density in this orientation is preferably 6 or more. From this viewpoint, the pole density is preferably 8 or more, more preferably 10 or more.
Furthermore, the pole density of {100} <011> in the plate thickness 1/2 layer ((001) [1-10] in the φ2 = 45 ° cross section of the ODF) significantly reduces the Young's modulus in the 45 ° direction. The pole density is desirably 6 or less. From this viewpoint, the pole density is preferably 3 or less, more preferably 1.5 or less.
Note that the crystal orientations described above allow variations within ± 2.5 °.
It is preferable that at least one of the above conditions for the crystal orientation pole density is satisfied, more preferably a plurality of conditions are satisfied, and the optimum condition is to satisfy all.
鋼板のヤング率については、上述した板厚１／８層と１／２層における結晶方位の極密度に関する要件を同時に満たすことで、圧延方向だけでなく、ＴＤ方向のヤング率も同時に２３０ＧＰａ超とすることが可能となる。ヤング率の測定はＪＩＳＺ２２８０に準拠した常温での横共振法にて行う。すなわち試料を固定せずに振動を加え、発振機の振動数を徐々に変化させて一次共振振動数を測定して、以下の［５］式よりヤング率を算出する。 Regarding the Young's modulus of the steel sheet, the Young's modulus not only in the rolling direction but also in the TD direction exceeds 230 GPa at the same time by simultaneously satisfying the requirements regarding the pole density of the crystal orientation in the 1/8 layer and 1/2 layer thicknesses described above. It becomes possible to do. The Young's modulus is measured by a transverse resonance method at room temperature in accordance with JISZ2280. In other words, vibration is applied without fixing the sample, the frequency of the oscillator is gradually changed, the primary resonance frequency is measured, and the Young's modulus is calculated from the following equation [5].
ここで、Ｅ：動的ヤング率［Ｎ／ｍ２］、ｌ：試験片の長さ［ｍ］、ｈ：試験片の厚さ［ｍ］、ｍ：質量［ｋｇ］、ｗ：試験片の幅［ｍ］、ｆ：横共振法の一次共振振動数［ｓ−１］、である。 Here, E: dynamic Young's modulus [N / m 2 ], l: length of test piece [m], h: thickness of test piece [m], m: mass [kg], w: of test piece Width [m], f: primary resonance frequency [s −1 ] of the lateral resonance method.
Ｘ線回折用試料の作製は、例えば次のようにして行う。
鋼板を機械研磨や化学研磨などによって板厚方向に所定の位置まで研磨し、バフ研磨によって鏡面に仕上げた後、電解研磨や化学研磨によって歪みを除去すると同時に板厚１／８層又は１／２層が測定面となるように調整する。例えば，１／８層の場合は、鋼板の板厚をｔとしたとき、ｔ／８の厚み分の研磨量で鋼板表面を研磨して現れる研磨面を測定面とする。なお、正確に板厚１／８層や１／２層を測定面とすることは困難であるので、これら目標とする層を中心として板厚に対して±３％の範囲が測定面となるように試料を作製すればよい。また、鋼板の板厚中心層に偏析帯が認められる場合には、板厚の３／８〜５／８の範囲で偏析帯のない場所について測定すればよい。さらにＸ線測定が困難な場合には、ＥＢＳＰ（Ｅｌｅｃｔｒｏｎ Ｂａｃｋ Ｓｃａｔｔｅｒ ｄｉｆｆｒａｃｔｉｏｎ Ｐａｔｔｅｒｎ）法やＥＣＰ（Ｅｌｅｃｔｒｏｎ Ｃｈａｎｎｅｌｉｎｇ Ｐａｔｔｅｒｎ）法により統計的に十分な数の測定を行う。
また、上記の｛ｈｋｌ｝＜ｕｖｗ＞は、上述の方法でＸ線用試料を採取したとき、板面に垂直な結晶方位が＜ｈｋｌ＞で圧延方向に平行な方位が＜ｕｖｗ＞であることを意味する。
For example, the X-ray diffraction sample is manufactured as follows.
The steel plate is polished to a predetermined position in the plate thickness direction by mechanical polishing or chemical polishing, and finished to a mirror surface by buffing, and then the strain is removed by electrolytic polishing or chemical polishing, and at the same time a plate thickness of 1/8 layer or 1/2 Adjust the layer to be the measurement surface. For example, in the case of 1/8 layer, when the plate thickness of the steel plate is t, the polished surface that appears by polishing the steel plate surface with the polishing amount corresponding to the thickness of t / 8 is taken as the measurement surface. In addition, since it is difficult to accurately set the plate thickness 1/8 layer or 1/2 layer as the measurement surface, the measurement surface has a range of ± 3% with respect to the plate thickness centering on these target layers. A sample may be prepared as described above. Moreover, when a segregation band is recognized in the plate thickness center layer of the steel sheet, it may be measured in a place where there is no segregation band in the range of 3/8 to 5/8 of the plate thickness. Further, when X-ray measurement is difficult, a statistically sufficient number of measurements are performed by an EBSP (Electron Back Scattering Pattern Pattern) method or an ECP (Electron Channeling Pattern) method.
In addition, the above {hkl} <uvw> means that when an X-ray sample is collected by the above method, the crystal orientation perpendicular to the plate surface is <hkl> and the orientation parallel to the rolling direction is <uvw>. Means.
固溶Ｃ量は０．０００５〜０．００４％とすることが好ましい。これを含有する鋼板が部材として加工されると、常温でも歪時効を生じ、ヤング率が高くなる。たとえば自動車用途に使用した場合に、加工後塗装焼付処理を施すことで鋼板の降伏強度のみならずヤング率も増加する。固溶Ｃ量は、全Ｃ量からＦｅ、Ａｌ、Ｎｂ、Ｔｉ、Ｂなどの化合物として存在するＣ量（抽出残査の化学分析から定量）を差し引いた値から求めることもできる。また、内部摩擦法やＦＩＭ（Ｆｉｅｌｄ Ｉｏｎ Ｍｉｃｒｏｓｃｏｐｙ）によって求めても良い。固溶Ｃ量が０．０００５％未満では十分な効果を得ることができない。また、０．００４％を超えてもＢＨ性は飽和する傾向にあるので、これを上限とする。 The amount of solute C is preferably 0.0005 to 0.004%. When a steel sheet containing this is processed as a member, strain aging occurs at room temperature, and the Young's modulus increases. For example, when used in automobile applications, the Young's modulus is increased as well as the yield strength of the steel sheet by performing post-processing paint baking. The amount of solid solution C can also be obtained from a value obtained by subtracting the amount of C existing as a compound such as Fe, Al, Nb, Ti, B, etc. (quantitative determination from chemical analysis of the extraction residue) from the total amount of C. Further, it may be obtained by an internal friction method or FIM (Field Ion Microscopy). If the amount of solute C is less than 0.0005%, a sufficient effect cannot be obtained. Further, even if it exceeds 0.004%, the BH property tends to be saturated, so this is the upper limit.
板厚の表層から１／８層における圧延方向及び幅方向のヤング率：板厚の表層から１／８層における圧延方向のヤング率の下限値は２４０ＧＰａとすることが好ましい。この値が２４０ＧＰａ未満では十分な曲げ剛性向上効果が得られない。この観点から圧延方向の表層ヤング率の下限は２４５ＧＰａとすることが望ましい。更に望ましくは２５０ＧＰａである。上限値は特に規定しないが、３００ＧＰａ超にするためには他の合金元素を大量に添加する必要があり、加工性等の他の特性が劣化することから実質３００ＧＰａ以下となる。また、表層のヤング率が２４０ＧＰａを超えていても、その層の厚みが１／８厚未満では十分な形状凍結性向上効果が発揮されない。高ヤング率を有する層の厚みは厚いほど高い曲げ剛性が得られるのは言うまでもない。
なお、表層のヤング率の測定は表層から１／８以上の厚みで試験片を切り出し、前述の横振動法にて行う。板幅方向の表層ヤング率は特に規定しないが、表層ヤング率が高い方が幅方向の曲げ剛性が上がることは言うまでもない。また上述のような成分製法によって幅方向の表層ヤング率も圧延方向と同様に２４０ＧＰａを超える。
The Young's modulus in the rolling direction and the width direction from the surface layer of the plate thickness to the 1/8 layer: The lower limit value of the Young's modulus in the rolling direction from the surface layer of the plate thickness to the 1/8 layer is preferably 240 GPa. If this value is less than 240 GPa, a sufficient bending rigidity improvement effect cannot be obtained. From this viewpoint, the lower limit of the surface Young's modulus in the rolling direction is preferably 245 GPa. More desirably, it is 250 GPa. Although the upper limit is not particularly defined, in order to make it over 300 GPa, it is necessary to add a large amount of other alloy elements, and other characteristics such as workability are deteriorated, so that it becomes substantially 300 GPa or less. Even if the Young's modulus of the surface layer exceeds 240 GPa, a sufficient effect of improving the shape freezing property cannot be exhibited if the thickness of the layer is less than 1/8. It goes without saying that the higher the rigidity of the layer having a high Young's modulus, the higher the bending rigidity.
In addition, the measurement of the Young's modulus of a surface layer cuts out a test piece with thickness of 1/8 or more from a surface layer, and performs the above-mentioned lateral vibration method. The surface Young's modulus in the plate width direction is not particularly defined, but it goes without saying that the higher the surface Young's modulus, the higher the bending rigidity in the width direction. Moreover, the surface layer Young's modulus of the width direction exceeds 240 GPa similarly to a rolling direction by the above component manufacturing methods.
次に、本発明に係る低降伏比型高ヤング率鋼板の製造条件の限定理由について述べる。 熱間圧延に供するスラブは特に限定するものではない。すなわち、連続鋳造スラブや薄スラブキャスターなどで製造したものであればよい。また、鋳造後に直ちに熱間圧延を行う連続鋳造−直接圧延（ＣＣ−ＤＲ）のようなプロセスにも適合する。
熱延鋼板製造時に、集合組織を制御するために以下のように製造条件を限定する必要がある。
熱延加熱温度は１０００℃以上とする。これは、後述する熱延仕上温度をＡｒ３変態点以上とするために必要な温度である。熱間圧延を行う際には以下の［６］式で計算される有効ひずみ量ε＊が０．４以上かつ圧下率の合計が５０％以上となるようにする。このときの圧延ロールと鋼板との摩擦係数を０．２超とする。以上の条件は表層の剪断集合組織を発達せしめ、圧延方向のヤング率を高めるのに必須の条件である。
Next, the reason for limiting the production conditions of the low yield ratio type high Young's modulus steel sheet according to the present invention will be described. The slab used for hot rolling is not particularly limited. That is, what was manufactured with the continuous casting slab, the thin slab caster, etc. should just be used. It is also compatible with processes such as continuous casting-direct rolling (CC-DR) in which hot rolling is performed immediately after casting.
It is necessary to limit the production conditions as follows in order to control the texture during hot-rolled steel plate production.
Hot rolling heating temperature shall be 1000 degreeC or more. This is a temperature necessary for setting the hot-rolling finishing temperature described later to the Ar3 transformation point or higher. When performing hot rolling, the effective strain amount ε * calculated by the following equation [6] is 0.4 or more and the total reduction ratio is 50% or more. At this time, the friction coefficient between the rolling roll and the steel sheet is set to be more than 0.2. The above conditions are indispensable conditions for developing the shear texture of the surface layer and increasing the Young's modulus in the rolling direction.
ここで、ｎは仕上げ熱延の圧延スタンド数、εｊはｊ番目のスタンドで加えられたひずみ、εｎはｎ番目のスタンドで加えられたひずみ、ｔｉはｉ〜ｉ＋１番目のスタンド間の走行時間［ｓ］、τｉは気体常数Ｒ（＝１．９８７）とｉ番目のスタンドの圧延温度Ｔｉ［Ｋ］によって以下の［７］式で計算できる。 Here, n is the number of finishing hot rolling rolling stands, εj is the strain applied at the jth stand, εn is the strain applied at the nth stand, ti is the travel time between the i to i + 1th stands [ s] and τi can be calculated by the following equation [7] according to the gas constant R (= 1.987) and the rolling temperature Ti [K] of the i-th stand.
有効ひずみε＊は０．４以上とし、０．５以上が好ましく、０．６以上であればより好ましい。圧下率の合計は５０％以上とし、板厚１／２層における｛２１１｝＜０１１＞方位、｛３３２｝＜１１３＞方位のいずれか一方または双方の極密度を大きくするためには、７０％以上が好ましく、１００％以上であればより好ましい。
圧下率の合計ＲＴは、ｎパスの圧延の場合、１パス目〜ｎパス目までの各圧下率をＲ１［％］〜Ｒｎ［％］とすると、以下の[８]式で示される。
The effective strain ε * is 0.4 or more, preferably 0.5 or more, and more preferably 0.6 or more. The total rolling reduction is 50% or more, and 70% is required to increase the pole density in one or both of the {211} <011> orientation and {332} <113> orientation in the 1/2 layer thickness. The above is preferable, and 100% or more is more preferable.
The total rolling reduction RT is expressed by the following formula [8], where each rolling reduction from the first pass to the n-th pass is R1 [%] to Rn [%] in the case of n-pass rolling.
ただし、Ｒｎ＝｛（ｎ−１）パス後の板厚−ｎパス後の板厚｝／（ｎ−１）パス後の板厚×１００［％］で定義できる。 However, Rn = {plate thickness after (n-1) pass-plate thickness after n pass} / (n-1) plate thickness after pass × 100 [%].
熱延の仕上温度は９００℃以下とする。９００℃超では、圧延方向に好ましい剪断集合組織を板厚表層から板厚１／４層付近まで発達させることが困難であり、｛１１０｝＜２２３＞、｛１１０｝＜１１１＞の双方の極密度が低下する。また、仕上げ温度を９００℃以下で、より低温にすると、γの再結晶が抑制されて板厚１／２層における｛１００｝＜０１１＞方位の極密度が低下し、４５°方向のヤング率には好ましくない集合組織の発達を抑制することができる。なお、せん断集合組織が発達した部位の表層からの深さを大きくするためにも、仕上圧延を低温で行うことが好ましい。この観点から、熱延の仕上温度は、好ましくは８５０℃以下、更に好ましくは８００℃以下とする。
一方、熱延の仕上温度は、Ａｒ３変態点未満では、｛１１０｝＜２２３＞、｛１１０｝＜１１１＞の双方の極密度が低下し、圧延方向のヤング率にとって好ましくない｛１１０｝＜００１＞集合組織が発達することもあるため、Ａｒ３変態点以上とする。
熱間圧延を実施する際には圧延ロールの異周速率が１％以上の異周速圧延を少なくとも１パス以上施すと表層近傍での集合組織形成が促進されるため、異周速圧延を実施しない場合の本発明以上にヤング率が向上する。この観点から異周速率は１％以上とし、望ましくは異周速率５％以上、更に望ましくは異周速率１０％以上の異周速圧延を施すことが望ましい。異周速率及び異周速圧延パス数の上限は特に規定しないが、上記の理由からいずれも大きい方が大きなヤング率向上効果が得られることは言うまでもない。しかし、５０％以上の異周速率は現状困難であり、仕上熱延パスは通常８パス程度までである。
ここで本発明における異周速率とは、上下圧延ロールの周速差を低周速側ロールの周速で除した値を百分率で表示したものである。また、本発明の異周速圧延は、上下ロール周速のいずれが大きくてもヤング率向上効果に差はない。
The finishing temperature of hot rolling is 900 ° C. or less. Above 900 ° C., it is difficult to develop a preferred shear texture in the rolling direction from the plate thickness surface layer to the vicinity of the 1/4 layer thickness, and both poles of {110} <223> and {110} <111> Density decreases. Further, when the finishing temperature is 900 ° C. or lower and lower, the recrystallization of γ is suppressed, the pole density in the {100} <011> orientation in the 1/2 layer thickness decreases, and the Young's modulus in the 45 ° direction It is possible to suppress unfavorable texture development. In order to increase the depth from the surface layer of the site where the shear texture has developed, it is preferable to perform the finish rolling at a low temperature. From this viewpoint, the finishing temperature of hot rolling is preferably 850 ° C. or lower, more preferably 800 ° C. or lower.
On the other hand, when the finishing temperature of hot rolling is less than the Ar3 transformation point, the extreme densities of both {110} <223> and {110} <111> are lowered, which is not preferable for the Young's modulus in the rolling direction. {110} <001 > Because a texture may develop, the Ar3 transformation point or higher is set.
When carrying out hot rolling, forming the texture in the vicinity of the surface layer is promoted if at least one pass of different circumferential speed rolling with a different circumferential speed ratio of the rolling roll of 1% or more is performed. If not, the Young's modulus is improved more than the present invention. From this point of view, it is desirable that the different peripheral speed rate is 1% or more, preferably the different peripheral speed ratio is 5% or more, and more preferably, the different peripheral speed ratio is 10% or more. Although the upper limit of the different peripheral speed ratio and the number of different peripheral speed rolling passes is not particularly defined, it goes without saying that a larger Young's modulus can be obtained with a larger value for the above reasons. However, a different peripheral speed ratio of 50% or more is currently difficult, and the finishing hot rolling pass is usually up to about 8 passes.
Here, the different peripheral speed ratio in the present invention is a value obtained by dividing the peripheral speed difference between the upper and lower rolling rolls by the peripheral speed of the low peripheral speed roll in percentage. Further, the different peripheral speed rolling of the present invention has no difference in Young's modulus improvement effect regardless of the upper and lower roll peripheral speeds.
また、仕上熱延に使用する圧延機にロール径が７００ｍｍ以下のワークロールを一つ以上使用すると表層近傍での集合組織形成が促進されるため、使用しない場合の本発明以上にヤング率が向上することからロール径７００ｍｍ以下のワークロールを使用することが望ましい。この観点から、ワークロール径は７００ｍｍ以下とし、６００ｍｍ以下であることが望ましく、５００ｍｍ以下とすることが更に望ましい。ワークロール径の下限は特に規定しないが、３００ｍｍ以下になると通板制御が困難になる。小径ロールを使用するパス数の上限は特に規定しないが、前述のように仕上熱延パスは通常８パス程度までである。
熱延中に変態組織制御を行う際、すなわち熱延板の組織をマルテンサイトを含む複合組織とする際には、上記製造条件に加え仕上げ圧延後の冷却条件が重要となるため、熱延仕上げ後、空冷時間を制限して制御冷却を行う。仕上圧延後は、直ちに制御冷却することが好ましいが、設備上の制約によって、制御冷却の開始までは空冷されることがある。
仕上圧延後の空冷時間が３０ｓを超えると、マルテンサイトが５％未満となるため、空冷時間は３０ｓ以内とする。
制御冷却の冷却速度は５℃／ｓ以上とし、制御冷却の停止温度を２５〜３００℃の範囲内とする。これは、熱延仕上後の冷却速度が５℃／ｓ未満では、低降伏比となるための体積分率の下限である２％以上のマルテンサイトを得るのが困難だからである。一方、冷却速度の上限は、特に制限に意味は無いが製造上１５０℃／ｓ以上にする事は困難である。なお、冷却速度を５〜１５０℃／ｓとする制御冷却は、水冷、ミスト冷却によって行うことができる。
In addition, when one or more work rolls having a roll diameter of 700 mm or less are used in a rolling mill used for finishing hot rolling, texture formation near the surface layer is promoted, and thus the Young's modulus is improved over the present invention when not used. Therefore, it is desirable to use a work roll having a roll diameter of 700 mm or less. From this viewpoint, the work roll diameter is 700 mm or less, desirably 600 mm or less, and more desirably 500 mm or less. The lower limit of the work roll diameter is not particularly defined, but if it is 300 mm or less, the sheet passing control becomes difficult. Although the upper limit of the number of passes using the small-diameter roll is not particularly defined, as described above, the finish hot rolling pass is usually up to about 8 passes.
When controlling the transformation structure during hot rolling, that is, when making the structure of the hot rolled sheet into a composite structure containing martensite, the cooling conditions after finish rolling are important in addition to the above manufacturing conditions. Thereafter, controlled cooling is performed by limiting the air cooling time. Although it is preferable to perform control cooling immediately after finish rolling, air cooling may occur until the start of control cooling due to restrictions on equipment.
If the air cooling time after finish rolling exceeds 30 s, martensite is less than 5%, so the air cooling time is set within 30 s.
The cooling rate of the controlled cooling is 5 ° C./s or more, and the controlled cooling stop temperature is in the range of 25 to 300 ° C. This is because if the cooling rate after hot rolling is less than 5 ° C./s, it is difficult to obtain martensite of 2% or more, which is the lower limit of the volume fraction for achieving a low yield ratio. On the other hand, the upper limit of the cooling rate is not particularly limited, but it is difficult to make it 150 ° C./s or more in production. Note that controlled cooling with a cooling rate of 5 to 150 ° C./s can be performed by water cooling or mist cooling.
熱延後の巻き取り温度は、２５〜３００℃とする。これは巻き取り温度が３００℃超ではマルテンサイトを２％以上得られないからである。また、下限値を２５℃としたのは、製造上の制約により２５℃未満にする事は困難だからである。
熱延鋼板には、最高加熱温度を５００℃以上０．５×（Ａｃ１＋Ａｃ３）［℃］以下の範囲とする焼鈍を施してもよい。最高加熱温度に到達後、直ちに冷却しても良いが、鋼板の温度を均一にするには、１２０ｓ以上保持することが好ましく、１８００ｓ超の保持は生産性を損なう。鋼板の材質の均質性と生産性を両立するには、保持時間を３００ｓ以上６００ｓ以下とすることが更に好ましい。また、熱延鋼板には、必要に応じて酸洗、インライン又はオフラインによる圧下率１０％以下のスキンパスを施しても良い。
The coiling temperature after hot rolling is 25 to 300 ° C. This is because if the winding temperature exceeds 300 ° C., 2% or more of martensite cannot be obtained. The lower limit is set to 25 ° C. because it is difficult to make it lower than 25 ° C. due to manufacturing restrictions.
The hot-rolled steel sheet may be annealed with a maximum heating temperature in the range of 500 ° C. or higher and 0.5 × (Ac1 + Ac3) [° C.] or lower. Although it may be cooled immediately after reaching the maximum heating temperature, in order to make the temperature of the steel plate uniform, it is preferable to hold for 120 s or more, and holding for over 1800 s impairs productivity. In order to achieve both the homogeneity of the material of the steel sheet and the productivity, the holding time is further preferably set to 300 s or more and 600 s or less. The hot-rolled steel sheet may be subjected to a skin pass having a reduction rate of 10% or less by pickling, in-line or off-line as necessary.
このようにして製造した熱延鋼板を酸洗後、冷延し、焼鈍しても良い。この場合、冷延率を６０％超とすると、熱延鋼板に形成されたヤング率を高める集合組織が大きく変化し、圧延方向のヤング率が低下してしまうことがあるため、冷延率は６０％未満とすることが好ましい。一方、冷延率の下限値に特に意味は無いが、生産性を考慮すると１０％以上とすることが好ましい。
冷間圧延後の焼鈍は、最高加熱温度を５００℃以上０．５×（Ａｃ１＋Ａｃ３）［℃］以下の範囲とし、２５〜３８０℃の範囲内まで冷却することが好ましい。これによって圧延方向のヤング率の低下を抑制し、かつ、体積分率で２〜２５％のマルテンサイトを確保することができる。熱延鋼板の焼鈍と同様、最高加熱温度に到達後、直ちに冷却しても良いが、１２０ｓ以上１８００ｓ以下の保持が好ましく、３００ｓ以上６００ｓ以下の保持が更に好ましい。
焼鈍後の冷却速度は、１℃／ｓ未満かつ冷却停止温度が３８０℃超では、冷却中にパーライト変態が起こる恐れがあるため、冷却速度を１℃／ｓ以上、冷却停止温度を３８０℃以下とすることが好ましい。なお、製造上、冷却速度を１５０℃／ｓ超、冷却停止温度を２５℃未満とすることは困難である。なお、パーライト変態を抑制するためには、冷却速度を２℃／ｓ以上とすることが好ましい。また、焼鈍後、２５〜３８０℃の範囲内に保持しても良いが、１８００ｓを超えるとマルテンサイトを確保できないことがあるため、保持時間は１８００ｓ以内とすることが好ましい。
The hot-rolled steel sheet thus manufactured may be cold-rolled and annealed after pickling. In this case, if the cold rolling rate is more than 60%, the texture that increases the Young's modulus formed in the hot-rolled steel sheet may change significantly, and the Young's modulus in the rolling direction may decrease. It is preferable to be less than 60%. On the other hand, the lower limit value of the cold rolling rate is not particularly meaningful, but is preferably 10% or more in consideration of productivity.
In the annealing after cold rolling, the maximum heating temperature is preferably in the range of 500 ° C. or more and 0.5 × (Ac1 + Ac3) [° C.] or less, and is preferably cooled to the range of 25 to 380 ° C. As a result, a decrease in Young's modulus in the rolling direction can be suppressed, and 2 to 25% martensite can be secured in terms of volume fraction. Similar to the annealing of the hot-rolled steel sheet, it may be cooled immediately after reaching the maximum heating temperature, but it is preferably 120 to 1800 s, more preferably 300 to 600 s.
If the cooling rate after annealing is less than 1 ° C / s and the cooling stop temperature exceeds 380 ° C, pearlite transformation may occur during cooling, so the cooling rate is 1 ° C / s or more and the cooling stop temperature is 380 ° C or less. It is preferable that In production, it is difficult to make the cooling rate higher than 150 ° C./s and the cooling stop temperature lower than 25 ° C. In order to suppress pearlite transformation, the cooling rate is preferably 2 ° C./s or more. Further, after annealing, the temperature may be maintained within a range of 25 to 380 ° C. However, if it exceeds 1800 s, martensite may not be ensured, and therefore the retention time is preferably within 1800 s.
熱間圧延後の冷却で変態組織を制御せず、冷間圧延後の焼鈍、冷却時に変態組織を制御してもよい。この場合も、熱間圧延によって、圧延方向のヤング率を向上させる剪断集合組織を表層に発達させるため、スラブ加熱温度、圧延ロールと鋼板との摩擦係数、有効ひずみ量ε＊、圧下率の合計、熱間圧延の終了温度は上記の条件を満足する必要がある。仕上圧延後の冷却条件は特に規定しないが、巻き取り条件については、４００〜６００℃で巻き取るとヤング率が向上する場合があるので、この範囲で巻き取るのが好ましい。また、熱延鋼板には最高加熱温度を５００℃以上０．５×（Ａｃ１＋Ａｃ３）［℃］以下、保持時間を０分以上、好ましくは１２０ｓ以上１８００ｓ以下、更に好ましくは３００ｓ以上６００ｓ以下の範囲とする焼鈍を施してもよい。 The transformation structure may be controlled at the time of annealing and cooling after cold rolling without controlling the transformation structure by cooling after hot rolling. Also in this case, the hot rolling is performed to develop a shear texture in the surface layer that improves the Young's modulus in the rolling direction, so that the slab heating temperature, the friction coefficient between the rolling roll and the steel sheet, the effective strain amount ε * , and the total reduction ratio The end temperature of the hot rolling needs to satisfy the above conditions. Although the cooling conditions after finish rolling are not particularly specified, the winding conditions are preferably in this range because the Young's modulus may be improved when winding at 400 to 600 ° C. The hot-rolled steel sheet has a maximum heating temperature of 500 ° C. or more and 0.5 × (Ac1 + Ac3) [° C.] or less, a holding time of 0 minutes or more, preferably 120 s or more and 1800 s or less, more preferably 300 s or more and 600 s or less. Annealing may be performed.
熱延鋼板を酸洗後、冷延を施し、冷延終了後の熱処理工程にて変態組織を制御する場合、冷延の圧下率、焼鈍の最高加熱温度、焼鈍後の冷却における冷却速度を限定する。冷延の圧下率の下限は製造上の観点から１０％とし、上限はヤング率の観点から６０％未満とする。焼鈍は、最高加熱温度をＡｃ１変態点［℃］以上０．５×（Ａｃ１＋Ａｃ３）［℃］とし、１℃／ｓ以上の冷却速度にて３８０℃以下まで冷却する。
ここで重要なのは、熱延で得られた集合組織の大部分を残しつつ、その一部を、フェライト相からオーステナイト相に変態（以下、α→γ変態という。）させ、その後の冷却制御にて低降伏比となるために必要なマルテンサイト量を得る事にある。すなわちＡｃ１変態点未満の最高加熱温度では、α→γ変態が起こりにくい事からこれを下限とし、また０．５×（Ａｃ１＋Ａｃ３）［℃］を超える温度では、組織の大部分が逆変態γとなり、熱延時に制御した集合組織制御によるヤング率向上効果が得られなくなる恐れのある事から、これを上限とした。加熱速度は特に限定しないが、３〜７０℃／ｓの範囲とすることが望ましい。加熱速度が３℃／ｓ未満では加熱中に再結晶が進行し、ヤング率向上に有利な集合組織がくずれてしまう。７０℃／ｓ超としても特段材料特性は変化しないことからこの値を上限とするのが望ましい。
When hot-rolled steel sheet is pickled, cold-rolled, and the transformation structure is controlled in the heat treatment process after the end of cold-rolling, the rolling reduction of cold rolling, the maximum heating temperature of annealing, and the cooling rate in cooling after annealing are limited To do. The lower limit of the cold rolling reduction is 10% from the viewpoint of manufacturing, and the upper limit is less than 60% from the viewpoint of Young's modulus. In the annealing, the maximum heating temperature is set to Ac1 transformation point [° C.] or more and 0.5 × (Ac1 + Ac3) [° C.], and it is cooled to 380 ° C. or less at a cooling rate of 1 ° C./s or more.
What is important here is that while most of the texture obtained by hot rolling remains, a part of the texture is transformed from the ferrite phase to the austenite phase (hereinafter referred to as α → γ transformation), and then the cooling control is performed. The purpose is to obtain the amount of martensite necessary to achieve a low yield ratio. That is, at the maximum heating temperature below the Ac1 transformation point, the α → γ transformation is unlikely to occur, so this is the lower limit. At temperatures exceeding 0.5 × (Ac1 + Ac3) [° C.], the majority of the tissue becomes the reverse transformation γ. Since the Young's modulus improvement effect by texture control controlled during hot rolling may not be obtained, this is set as the upper limit. The heating rate is not particularly limited, but is preferably in the range of 3 to 70 ° C./s. If the heating rate is less than 3 ° C./s, recrystallization proceeds during heating, and the texture advantageous for improving the Young's modulus is broken. Even if it exceeds 70 ° C./s, the material characteristics do not change, so it is desirable to set this value as the upper limit.
また、その後の冷却速度が１℃／ｓ未満かつ冷却停止温度が３８０℃超では、冷却中にパーライト変態が起こる恐れがあるため、これを下限と及び上限とした。一方、冷却速度の制限に意味は無いが製造上１５０℃／ｓ超にする事は困難であるため１５０℃／ｓ以下とし、冷却停止温度も同様の理由で２５℃以上とする。なお、パーライト変態の抑制という観点から、冷却速度は２℃／ｓ以上であることが好ましい。上記の焼鈍後、２５〜３８０℃の温度で保持しても良いが、１８００ｓを超えるとマルテンサイトを確保できないことがあるため、保持時間は１８００ｓ以内とすることが好ましい。
冷延鋼板には、必要に応じて酸洗した後、さらにインライン又はオフラインで圧下率１０％以下のスキンパスを施しても良い。
In addition, if the subsequent cooling rate is less than 1 ° C./s and the cooling stop temperature exceeds 380 ° C., pearlite transformation may occur during cooling, so this was set as the lower limit and the upper limit. On the other hand, although there is no meaning in limiting the cooling rate, it is difficult to make it higher than 150 ° C./s because of manufacturing, so that it is 150 ° C./s or lower, and the cooling stop temperature is 25 ° C. or higher for the same reason. From the viewpoint of suppressing pearlite transformation, the cooling rate is preferably 2 ° C./s or more. After the above annealing, it may be held at a temperature of 25 to 380 ° C. However, if it exceeds 1800 s, martensite may not be ensured, so the holding time is preferably set to 1800 s or less.
The cold-rolled steel sheet may be pickled as necessary, and then further subjected to in-line or off-line skin pass with a rolling reduction of 10% or less.
熱延鋼板、冷延鋼板には溶融亜鉛めっき又は合金化溶融亜鉛めっきを施してもよい。冷延鋼板を焼鈍する場合は、冷却後、連続する溶融亜鉛めっきラインにて、そのまま溶融亜鉛めっきを施してもよい。亜鉛めっきの組成は特に限定するものではなく、亜鉛のほか、Ｆｅ、Ａｌ、Ｍｎ、Ｃｒ、Ｍｇ、Ｐｂ、Ｓｎ、Ｎｉなどを必要に応じて添加しても構わない。
合金化処理は、溶融亜鉛めっきを施した後に、４５０〜６００℃の範囲内で行う。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、めっき層が脆化するため、プレス等の加工によってめっきが剥離するなどの問題を誘発する。合金化処理の時間は、５ｓ以上とする。５ｓ未満では合金化が十分に進行しない。上限は特に定めないが、めっき密着性を考慮すると１０ｓ程度とすることが好ましい。
焼鈍及び溶融亜鉛めっき後に、必要に応じて酸洗し、その後インライン又はオフラインで圧下率１０％以下のスキンパスを施しても良い。
冷延後に連続溶融亜鉛めっきラインにて熱処理及び亜鉛めっきを行っても構わない。
また、上記の熱延鋼板、冷延鋼板にはＡｌ系めっきや各種電気めっきを施しても構わない。さらに熱延鋼板や冷延鋼板及び各種めっき鋼板には有機皮膜、無機皮膜、各種塗料などの表面処理を目的に応じて行うことができる。
Hot-rolled steel sheets and cold-rolled steel sheets may be hot dip galvanized or alloyed hot dip galvanized. When the cold-rolled steel sheet is annealed, it may be subjected to hot dip galvanization as it is in a continuous hot dip galvanizing line after cooling. The composition of the galvanizing is not particularly limited, and besides zinc, Fe, Al, Mn, Cr, Mg, Pb, Sn, Ni, etc. may be added as necessary.
The alloying treatment is performed within a range of 450 to 600 ° C. after hot dip galvanizing. If it is less than 450 ° C, alloying does not proceed sufficiently, and if it exceeds 600 ° C, alloying proceeds excessively and the plating layer becomes brittle, which causes problems such as peeling of the plating by processing such as pressing. . The alloying time is 5 s or longer. If it is less than 5 s, alloying does not proceed sufficiently. Although the upper limit is not particularly defined, it is preferably about 10 s in consideration of plating adhesion.
After annealing and hot dip galvanizing, pickling may be performed as necessary, and then a skin pass with a rolling reduction of 10% or less may be applied in-line or offline.
Heat treatment and galvanization may be performed in a continuous hot dip galvanizing line after cold rolling.
Moreover, you may give Al type plating and various electroplating to said hot-rolled steel plate and cold-rolled steel plate. Furthermore, surface treatments such as organic coatings, inorganic coatings, and various paints can be applied to hot-rolled steel sheets, cold-rolled steel sheets, and various plated steel sheets depending on the purpose.
本発明の低降伏比型高ヤング率鋼板、溶融亜鉛めっき鋼板、合金化溶融亜鉛めっき鋼板を圧延方向が鋼管の長手方向との間の角度が０〜３０°以内になるように巻いて鋼管にすると、鋼管の長手方向のヤング率が高い高ヤング率鋼管を製造することができる。圧延方向と平行に巻くのが最もヤング率が高くなることからこの角度は出来るだけ小さいことが好ましい。この観点から、１５°以下の角度で巻くことが更に好ましい。圧延方向と鋼管の長手方向の関係が満足されていれば、造管方法はＵＯ管、電縫溶接、スパイラル等、任意の方法をとることができる。もちろん、ヤング率の高い方向を鋼管の長手方向に平行に限定する必要はなく、用途に応じて任意の方向にヤング率の高い鋼管を製造しても何ら問題はない。 The low yield ratio type high Young's modulus steel sheet, hot-dip galvanized steel sheet, and alloyed hot-dip galvanized steel sheet of the present invention are wound on a steel pipe so that the angle between the rolling direction and the longitudinal direction of the steel pipe is within 0 to 30 °. Then, a high Young's modulus steel pipe having a high Young's modulus in the longitudinal direction of the steel pipe can be manufactured. It is preferable that this angle be as small as possible since the Young's modulus is the highest when it is wound parallel to the rolling direction. From this viewpoint, it is more preferable to wind at an angle of 15 ° or less. As long as the relationship between the rolling direction and the longitudinal direction of the steel pipe is satisfied, the pipe forming method may be any method such as UO pipe, electric resistance welding, spiral, or the like. Of course, it is not necessary to limit the direction with a high Young's modulus parallel to the longitudinal direction of the steel pipe, and there is no problem even if a steel pipe with a high Young's modulus is produced in any direction depending on the application.
次に本発明を実施例にて説明する。
「実施例１」
表１に示す組成を有する鋼を溶製し、表２に示す条件で熱間圧延を施した。Ａｒ３、Ａｃ１及びＡｃ３は、フォーマスター試験機を用い、冷却中の試験片の熱膨張変化を測定して求めた。表１の成分の空欄は分析値が検出限界未満であることを意味し、Ｔｉ−４８／１４Ｎが空欄であるものはＴｉを含有しないものである。Ｍｏ＋Ｎｂ＋Ｂ＋Ｔｉは、各元素の含有量が検出限界未満であるものは０として計算した。Ｍｏ＋Ｎｂ＋Ｂ＋Ｔｉが０であるものは、Ｍｏ、Ｎｂ、Ｂ、Ｔｉの全ての含有量が検出限界未満であることを意味する。なお、表１及び２の下線は、また表３〜９においても同様に、本発明の範囲外又は好ましい条件の範囲外であることを意味する。
熱間圧延の加熱温度は全て１２３０℃とした。全７段からなる仕上圧延スタンドにおいてロールと鋼板との摩擦係数を０．２１〜０．２４の範囲とし、最終３段の合計の圧下率を５５％とした。調質圧延圧下率はすべて０．３％とした。光学顕微鏡による組織観察と画像解析によって、マルテンサイト体積率、フェライト体積率、ベイナイト体積率を求めた。なお、マルテンサイト、フェライト、ベイナイトの残部はパーライトである。
ヤング率は上述した横共振法により測定した。Ｅ（ＲＤ）、Ｅ（Ｄ）、Ｅ（ＴＤ）は、それぞれ、長手方向をＲＤ方向、４５°方向、ＴＤ方向として試験片を採取し、測定して得られた室温におけるヤング率である。ＪＩＳ Ｚ ２２０１の５号引張試験片を採取してＴＤ方向の引張特性をＪＩＳ Ｚ ２２４１に準拠して評価した。また、板厚１／８層及び板厚７／１６層における集合組織をＸ線回折法によって測定した。表２に熱延条件と熱延鋼板の集合組織を示し、表３に熱延鋼板の組織、引張特性及びヤング率を示す。
なお、表２において、ＦＴ［℃］は熱間圧延機で仕上圧延を行う際の最終スタンドの出側の温度であり、ＣＴ［℃］は巻き取り温度であり、ｔＡＣ［ｓ］は熱間圧延における仕上圧延後、冷却を開始するまでの空冷時間であり、ＣＲ１［℃/ｓ］はＦＴからＣＴまでの平均冷却速度である。また、表３は熱延鋼板の組織及び特性を示すものであり、ＶＭ１［％］はマルテンサイト体積率であり、Ｖα１はフェライト体積率であり、ＶＢ１［％］はベイナイト体積率である。ＴＳ［ＭＰａ］は引張強さであり、ＹＳ［ＭＰａ］は降伏強さであり、ＹＲ［−］は降伏比であり、Ｅｌ［％］は伸びであり、Ｅ（ＲＤ）［ＧＰａ］はＲＤ方向の平均ヤング率であり、Ｒ（Ｄ）［ＧＰａ］はＲＤ方向に対して４５°傾斜した方向の平均ヤング率であり、Ｅ（ＴＤ）［ＧＰａ］はＴＤ方向の平均ヤング率である。これらの指標は、以後の表４〜９の説明において共通する。
Next, the present invention will be described with reference to examples.
"Example 1"
Steel having the composition shown in Table 1 was melted and hot rolled under the conditions shown in Table 2. Ar3, Ac1 and Ac3 were determined by measuring the thermal expansion change of the test piece during cooling using a Formaster tester. The blank of the component in Table 1 means that the analytical value is less than the detection limit, and Ti-48 / 14N is blank and does not contain Ti. Mo + Nb + B + Ti was calculated assuming that the content of each element was less than the detection limit. When Mo + Nb + B + Ti is 0, it means that the contents of all of Mo, Nb, B, and Ti are less than the detection limit. In addition, the underline of Table 1 and 2 also means outside the range of this invention or the range of preferable conditions also in Tables 3-9.
The heating temperature of hot rolling was all set to 1230 ° C. In the finishing rolling stand consisting of all seven stages, the friction coefficient between the roll and the steel sheet was in the range of 0.21 to 0.24, and the total rolling reduction of the last three stages was 55%. All the temper rolling reduction ratios were 0.3%. The martensite volume fraction, ferrite volume fraction, and bainite volume fraction were determined by structural observation and image analysis using an optical microscope. The balance of martensite, ferrite, and bainite is pearlite.
Young's modulus was measured by the transverse resonance method described above. E (RD), E (D), and E (TD) are Young's moduli at room temperature obtained by collecting test specimens with the longitudinal direction as the RD direction, 45 ° direction, and TD direction, respectively. A No. 5 tensile test piece of JIS Z 2201 was collected and evaluated for tensile properties in the TD direction in accordance with JIS Z 2241. Moreover, the texture in the plate thickness 1/8 layer and the plate thickness 7/16 layer was measured by the X-ray diffraction method. Table 2 shows the hot rolling conditions and the texture of the hot rolled steel sheet, and Table 3 shows the structure, tensile properties, and Young's modulus of the hot rolled steel sheet.
In Table 2, FT [° C.] is the temperature on the exit side of the final stand when finish rolling is performed by a hot rolling mill, CT [° C.] is the winding temperature, and t AC [s] is the heat It is the air cooling time until the cooling is started after the finish rolling in the intermediate rolling, and CR 1 [° C./s] is an average cooling rate from FT to CT. Table 3 shows the structure and properties of the hot-rolled steel sheet, where V M1 [%] is the martensite volume fraction, Vα 1 is the ferrite volume fraction, and V B1 [%] is the bainite volume fraction. is there. TS [MPa] is the tensile strength, YS [MPa] is the yield strength, YR [−] is the yield ratio, El [%] is the elongation, and E (RD) [GPa] is the RD R (D) [GPa] is an average Young's modulus in a direction inclined by 45 ° with respect to the RD direction, and E (TD) [GPa] is an average Young's modulus in the TD direction. These indices are common in the following description of Tables 4-9.
表２及び３から明らかなとおり、本発明の化学成分を有する鋼を適正な条件で熱間圧延した場合には、圧延方向のヤング率を２３０ＧＰａ超とすることができた。 As apparent from Tables 2 and 3, when the steel having the chemical component of the present invention was hot-rolled under appropriate conditions, the Young's modulus in the rolling direction could be over 230 GPa.
「実施例２」
表１における鋼Ｎｏ．Ｑ、Ｒ及びＳの組成を有する鋼スラブを溶製し、表４に示す条件で熱間圧延を施した。スラブの加熱温度、仕上圧延スタンド数、ロールと鋼板との摩擦係数、最終３段の合計の圧下率は実施例１と同様にして行った。また、Ａｒ３は表２の場合と同様に示し、更にＡｃ１及びＡｃ３も併記した。
このようにして得られた熱延鋼板を酸洗し、表４に示す圧下率で冷間圧延後、連続焼鈍（最高加熱温度にて９０ｓ保持）、連続溶融亜鉛めっき（最高加熱温度にて９０ｓ保持後、亜鉛めっき浴に浸漬後５００℃で１０ｓの合金化処理を実施）の何れかの処理を施し、引張特性とヤング率を測定した。マルテンサイト体積率、フェライト体積率、ベイナイト体積率、集合組織、ヤング率の測定及び引張特性の評価は実施例１と同様にして行った。調質圧延の圧下率はすべて０．３％とした。結果を表５に示す。
なお、表４において、ＣＲ２［℃/ｓ］は焼鈍後の平均冷却速度であり、ＴＯＡ［℃］は焼鈍中の冷却後の保持温度であり、tＯＡ［ｓ］は焼鈍中の冷却後の保持時間である。また、表５において、ＶＭ２は冷間圧延焼鈍後のマルテンサイト体積率であり、Ｖα２は冷間圧延焼鈍後のフェライト体積率であり、ＶＢ２は冷間圧延焼鈍後のベイナイト体積率である。
"Example 2"
Steel No. 1 in Table 1 Steel slabs having compositions of Q, R and S were melted and hot rolled under the conditions shown in Table 4. The heating temperature of the slab, the number of finish rolling stands, the coefficient of friction between the roll and the steel sheet, and the total rolling reduction of the final three stages were the same as in Example 1. Ar3 is shown in the same manner as in Table 2, and Ac1 and Ac3 are also shown.
The hot-rolled steel sheet thus obtained was pickled and cold-rolled at the rolling reduction shown in Table 4, followed by continuous annealing (maintained at the maximum heating temperature for 90 s) and continuous hot dip galvanizing (at the maximum heating temperature of 90 s). After holding, the steel sheet was immersed in a galvanizing bath and then subjected to an alloying treatment at 500 ° C. for 10 s), and tensile properties and Young's modulus were measured. Measurement of martensite volume fraction, ferrite volume fraction, bainite volume fraction, texture, Young's modulus, and evaluation of tensile properties were carried out in the same manner as in Example 1. The rolling reduction of temper rolling was all 0.3%. The results are shown in Table 5.
In Table 4, CR 2 [° C./s] is an average cooling rate after annealing, T OA [° C.] is a holding temperature after cooling during annealing, and t OA [s] is cooling during annealing. Later holding time. In Table 5, V M2 is the volume fraction of martensite after cold rolling annealing, V.alpha 2 is a ferrite volume fraction after cold rolling annealing, V B2 at bainite volume fraction after cold rolling annealing is there.
表５に示す試験結果から明らかなとおり、本発明の化学成分を有する鋼を適正な条件で熱延し、冷間圧延した後、さらに適切に熱処理することによって、冷延鋼板及びめっき鋼板の圧延方向のヤング率が向上する。 As is clear from the test results shown in Table 5, the steel having the chemical components of the present invention is hot-rolled under appropriate conditions, cold-rolled, and then further appropriately heat-treated, thereby rolling the cold-rolled steel sheet and the plated steel sheet. The Young's modulus in the direction is improved.
「実施例３」
表１に示した鋼ＢとＯを用いて異周速圧延を行った。周速率は全７段からなる仕上げ圧延スタンドにおいて最終の３段で変化させた。マルテンサイト体積率、フェライト体積率、ベイナイト体積率、集合組織、ヤング率の測定及び引張特性の評価は実施例１と同様にして行った。熱延条件及び集合組織を表６に、各組織の体積率、引張特性とヤング率の測定結果を表７に示す。なお、表６で表示されていない熱延条件は全て実施例１と同様である。
表６及び７に示す試験結果から明らかなとおり、本発明の化学成分を有する鋼を適正な条件で熱延する際に１％以上の異周速圧延を１パス以上加えると、表層近傍での集合組織形成が促進され、更にヤング率が向上する。
"Example 3"
Different peripheral speed rolling was performed using steel B and O shown in Table 1. The peripheral speed ratio was changed in the final three stages in a finish rolling stand having a total of seven stages. Measurement of martensite volume fraction, ferrite volume fraction, bainite volume fraction, texture, Young's modulus, and evaluation of tensile properties were carried out in the same manner as in Example 1. Table 6 shows the hot rolling conditions and texture, and Table 7 shows the measurement results of the volume ratio, tensile properties, and Young's modulus of each structure. All the hot rolling conditions not displayed in Table 6 are the same as in Example 1.
As is apparent from the test results shown in Tables 6 and 7, when hot rolling the steel having the chemical composition of the present invention under appropriate conditions, adding 1% or more of different peripheral speed rolling more than 1 pass, Texture formation is promoted and Young's modulus is further improved.
「実施例４」
表１に示した鋼ＢとＯを用いて異周速圧延を行った。周速率は全７段からなる仕上げ圧延スタンドにおいて最終の３段で変化させた。マルテンサイト体積率、フェライト体積率、ベイナイト体積率、集合組織、ヤング率の測定及び引張特性の評価は実施例１と同様にして行った。熱延条件及び集合組織を表８に、各組織の体積率、引張特性とヤング率の測定結果を表９に示す。なお、表８で表示されていない熱延条件は全て実施例１と同様である。
Example 4
Different peripheral speed rolling was performed using steel B and O shown in Table 1. The peripheral speed ratio was changed in the final three stages in a finish rolling stand having a total of seven stages. Measurement of martensite volume fraction, ferrite volume fraction, bainite volume fraction, texture, Young's modulus, and evaluation of tensile properties were carried out in the same manner as in Example 1. Table 8 shows the hot rolling conditions and texture, and Table 9 shows the measurement results of the volume ratio, tensile properties, and Young's modulus of each structure. The hot rolling conditions not displayed in Table 8 are all the same as in Example 1.
表８及び９に示す試験結果から明らかなとおり、本発明の化学成分を有する鋼を適正な条件で熱延する際に１％以上の異周速圧延を１パス以上加えると、表層近傍での集合組織形成が促進され、更にヤング率が向上する。
As is clear from the test results shown in Tables 8 and 9, when hot rolling the steel having the chemical components of the present invention under appropriate conditions, adding 1% or more of different peripheral speed rolling more than one pass, Texture formation is promoted and Young's modulus is further improved.
Claims (22)
Ｃ：０．０００５〜０．３０％、Ｓｉ：２．５％以下、
Ｍｎ：０．１〜５．０％、Ｐ：０．１５％以下、
Ｓ：０．０１５％以下、Ａｌ：０．１５％以下
Ｎ：０．０１％以下、
及びＭｏ：０．００５〜１．５％、Ｎｂ：０．００５〜０．２０％ 、Ｔｉ：４８／１４×Ｎ［ｍａｓｓ％］以上、０．２％以下、Ｂ：０．０００１〜０．０１％、のいずれか１種又は２種以上を合計で０．０１５〜１．９１％含有し、残部鉄及び不可避的不純物からなり、フェライト又はベイナイトを体積分率最大の組織とし、体積分率で２〜２５％のマルテンサイトを含む複合組織鋼であり、かつ板厚の１／８層における｛１１０｝＜２２３＞、｛１１０｝＜１１１＞の一方又は双方の極密度が１０以上であり、圧延方向のヤング率が２３０ＧＰａ超であることを特徴とする低降伏比型高ヤング率鋼板。 % By mass
C: 0.0005 to 0.30%, Si: 2.5% or less,
Mn: 0.1 to 5.0%, P: 0.15% or less,
S: 0.015% or less, Al: 0.15% or less N: 0.01% or less,
And Mo: 0.005-1.5%, Nb: 0.005-0.20%, Ti: 48/14 × N [mass%] or more, 0.2% or less, B: 0.0001-0. 01%, or a total of 0.015 to 1.91% of any one or two of them, consisting of the balance iron and inevitable impurities, with ferrite or bainite as the largest volume fraction, and volume fraction And the pole density of one or both of {110} <223> and {110} <111> in the 1/8 layer of the plate thickness is 10 or more. A low yield ratio type high Young's modulus steel sheet having a Young's modulus in the rolling direction of more than 230 GPa.
Ｂ：０．０００６〜０．０１％、
Ｎｂ：０．０１〜０．２０％、
Ｔｉ：４８／１４×Ｎ［ｍａｓｓ％］以上、０．２％以下
を含有し、板厚の１／８層における｛１１０｝＜００１＞の極密度が６以下であることを特徴とする請求項１記載の低降伏比型高ヤング率鋼板。 Mo: 0.1 to 1.5%
B: 0.0006 to 0.01%
Nb: 0.01-0.20%,
Ti: 48/14 × N [mass%] or more and 0.2% or less, and the {110} <001> pole density in the 1/8 layer of the plate thickness is 6 or less. Item 2. A low yield ratio type high Young's modulus steel sheet according to Item 1.
A method for producing a steel pipe according to claim 13, wherein the steel plate is a low yield ratio type high Young's modulus steel plate obtained by the production method according to any one of claims 14 to 18, and the production according to claim 19 or 20. A low yield ratio type high Young's modulus hot-dip galvanized steel sheet obtained by the method or a low yield ratio type high Young's modulus alloyed hot-dip galvanized steel sheet obtained by the manufacturing method according to claim 21 is wound in an arbitrary direction into a steel pipe. A method for producing a low yield ratio type high Young's modulus steel pipe.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2006116797A JP5058508B2 (en) | 2005-11-01 | 2006-04-20 | Low yield ratio type high Young's modulus steel plate, hot dip galvanized steel plate, alloyed hot dip galvanized steel plate and steel pipe, and production method thereof |
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| Application Number | Priority Date | Filing Date | Title |
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