JP2004218077A - Good burring property high strength steel sheet excellent in softening resistance in welded heat affecting zone, and its production method - Google Patents
Good burring property high strength steel sheet excellent in softening resistance in welded heat affecting zone, and its production method Download PDFInfo
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Abstract
Description
本発明は、溶接熱影響部の耐軟化性に優れた引張強度５４０ＭＰａ以上のバーリング性高強度鋼板およびその製造方法に関するものであり、特に、成形後にスポット、アーク、プラズマ、レーザー等により溶接される場合や、これら溶接後に成形される場合において加工性と溶接部の強度の両立が求められる自動車部品等の用途に用いられる素材として好適な溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板およびその製造方法に関するものである。 The present invention relates to a high-strength burring steel sheet having a tensile strength of 540 MPa or more and excellent in softening resistance of a weld heat-affected zone, and a method for producing the same. In particular, it is welded by spot, arc, plasma, laser or the like after forming. High strength burring with excellent softening resistance of welding heat affected zone, suitable as a material used in applications such as automotive parts where both workability and strength of the welded part are required when forming after welding. The present invention relates to a steel plate and a method for manufacturing the same.
近年、自動車の燃費向上などのために軽量化を目的として、Ａｌ合金等の軽金属や高強度鋼板の自動車部材への適用が進められている。 BACKGROUND ART In recent years, application of light metals such as Al alloys and high-strength steel sheets to automobile members has been promoted for the purpose of weight reduction in order to improve fuel efficiency of automobiles.
ただ、Ａｌ合金等の軽金属は比強度が高いという利点があるものの鋼に比較して著しく高価であるためその適用は特殊な用途に限られてきた。より広い範囲で自動車の軽量化を推進するためには安価な高強度鋼板の適用が強く求められている。 However, although light metals such as Al alloys have the advantage of high specific strength, their application has been limited to special applications because they are significantly more expensive than steel. In order to reduce the weight of automobiles over a wider range, there is a strong demand for the use of inexpensive high-strength steel sheets.
一般に材料は高強度になるほど成形性が悪くなる。鉄鋼材料においても例外ではなく、これまでに高強度と高延性の両立の試みがなされてきた。また、自動車部品に使用される材料に求められる特性としては延性の他にバーリング加工性がある。しかし、バーリング加工性も高強度化に伴って低下する傾向を示すため、バーリング加工性の向上も高強度鋼板の自動車部品への適用の課題となっている。一方、自動車部品はプレス成形等によって加工された部材をスポット、アーク、プラズマ、レーザー等の溶接によってアッセンブルされる。また、最近では鋼板をこれら溶接によって接合した後にプレス成形される場合もある。いずれにしても成形時もしくは部品として組み付けられて使用された時の溶接部強度は成形限界、安全性の面から非常に重要である。従って、自動車部品等への高強度鋼板の適用にあたっては、そのバーリング加工性とともに溶接部強度も重要な検討課題となる。 Generally, the higher the strength of a material, the worse the formability. Steel materials are no exception, and attempts have been made to achieve both high strength and high ductility. Properties required for materials used for automobile parts include burring workability in addition to ductility. However, since the burring workability also tends to decrease with the increase in strength, improvement of the burring workability is also an issue of applying high-strength steel sheets to automobile parts. On the other hand, automotive parts are assembled by welding spots, arcs, plasmas, lasers, or the like to members processed by press molding or the like. Recently, there is also a case where a steel sheet is press-formed after being joined by these weldings. In any case, the strength of the weld at the time of molding or when assembled and used as a part is very important from the viewpoint of molding limitations and safety. Therefore, when applying a high-strength steel sheet to an automobile part or the like, not only the burring workability but also the strength of the welded portion are important considerations.
バーリング加工性に優れた高強度鋼板については、Ｔｉ、Ｎｂを添加することにより第二相を低減し主相であるポリゴナルフェライト中にＴｉＣ、ＮｂＣを析出強化させることによって伸びフランジ性の優れた高強度熱延鋼板とした発明がある（例えば、特許文献１）。 For high-strength steel sheets excellent in burring workability, by adding Ti and Nb, the second phase is reduced, and TiC and NbC are precipitated and strengthened in the polygonal ferrite, which is the main phase, so that the stretch flangeability is excellent. There is an invention using a high-strength hot-rolled steel sheet (for example, Patent Document 1).
また、Ｔｉ、Ｎｂを添加することにより第二相を低減しミクロ組織をアシキュラーフェライトとしＴｉＣ、ＮｂＣで析出強化することによって伸びフランジ性の優れた高強度熱延鋼板とした発明がある（例えば、特許文献２）。 In addition, there is an invention in which a high-strength hot-rolled steel sheet having excellent stretch flangeability is obtained by adding Ti and Nb to reduce the second phase, change the microstructure to acicular ferrite, and strengthen the precipitation with TiC and NbC (for example, , Patent Document 2).
一方、溶接部強度を改善する技術としては、Ｎｂ、Ｍｏの複合添加により溶接部の軟化を抑制する鋼板を得る発明がある（例えば、特許文献３）。 On the other hand, as a technique for improving the strength of a weld, there is an invention for obtaining a steel sheet that suppresses the softening of the weld by a combined addition of Nb and Mo (for example, Patent Document 3).
また、ＮｂＮの析出を活用して溶接部の軟化を抑制するフェライトおよびマルテンサイトからなる鋼板を得る発明もある（例えば、特許文献４）。 In addition, there is also an invention for obtaining a steel sheet made of ferrite and martensite, which suppresses softening of a welded portion by utilizing precipitation of NbN (for example, Patent Document 4).
しかしながら、サスペンションアームやフロントサイドメンバー等一部の部品用鋼板においては、バーリング加工性をはじめとする成形性とともに溶接部の強度が大変に重要であり、上記従来技術では、これら両特性を共に満足することができない。また例え両特性が満足されたとしても安価に安定して製造できる製造方法を提供することが重要であり、上記従来技術では、不十分であると言わざるを得ない。 However, in the case of steel plates for some parts such as suspension arms and front side members, the formability such as burring workability and the strength of the welded part are very important, and the above conventional technology satisfies both these characteristics. Can not do it. Even if both characteristics are satisfied, it is important to provide a manufacturing method that can be manufactured stably at a low cost, and it cannot be said that the above conventional technology is insufficient.
すなわち、特許文献１に記載の発明では、高い伸びフランジ性を得るために面積率で８５％以上のポリゴナルフェライトが必須であるが、８５％以上のポリゴナルフェライトを得るためには熱間圧延後にフェライト粒の成長を促進するため長時間の保持が必要であり操業コスト上好ましくない。 That is, in the invention described in Patent Literature 1, polygonal ferrite having an area ratio of 85% or more is indispensable to obtain high stretch flangeability, but hot rolling is performed to obtain polygonal ferrite of 85% or more. Long-term holding is necessary to promote ferrite grain growth later, which is not preferable in terms of operating costs.
また、特許文献２に記載の発明では、転位密度が高いミクロ組織と微細なＴｉＣ及び／又はＮｂＣの析出によって８０ｋｇｆ／ｍｍ2で１７％程度の延性しかなく成形性が不十分である。
In the invention described in
さらに、これらの発明は溶接部の軟化については何ら言及していいない。一方、特許文献３に記載の発明には、バーリング加工性向上に関しては何も記載されていない。
Furthermore, these inventions do not mention softening of the weld. On the other hand, the invention described in
さらに、特許文献４に記載の発明は、フェライト−マルテンサイト複合組織鋼に関するものでは本発明のバーリング加工性に優れる鋼板のミクロ組織を得る技術とは明らかに異なる。
Furthermore, the invention described in
本発明は、上記問題点を解決して成形後にスポット、アーク、プラズマ、レーザー等により溶接される場合や、これら溶接後に成形される場合において加工性と溶接部の強度の両立が求められる自動車部品等の用途に用いられる素材として好適な溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板およびその製造方法を得ようとするものである。すなわち、本発明は、溶接熱影響部の耐軟化性に優れた引張強度５４０ＭＰａ以上のバーリング性高強度鋼板およびその鋼板を安価に安定して製造できる製造方法を提供することを目的とするものである。 The present invention solves the above-described problems, and spots, arcs, plasmas, lasers, and the like are welded after molding, and automobile parts that require both workability and strength of a welded part when molded after these weldings are required. It is an object of the present invention to obtain a burring high-strength steel sheet excellent in softening resistance of a weld heat-affected zone suitable as a material used for applications such as the above, and a method for producing the same. That is, an object of the present invention is to provide a high-strength burring steel sheet having a tensile strength of 540 MPa or more excellent in softening resistance of a weld heat-affected zone, and a manufacturing method capable of stably manufacturing the steel sheet at low cost. is there.
本発明者らは、現在通常に採用されている製造設備により工業的規模で生産されている薄鋼板の製造プロセスを念頭において、バーリング性高強度鋼板の溶接熱影響部の耐軟化性を向上させるべく鋭意研究を重ねた。その結果、Ｃ：０．０１〜０．１％、Ｓｉ：０．０１〜２％、Ｍｎ：０．０５〜３％、Ｐ≦０．１％、Ｓ≦０．０３％、Ａｌ：０．００５〜１％、Ｎ：０．０００５〜０．００５％、Ｔｉ：０．０５〜０．５％、を含み、さらに０＜Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％、さらに、Ｍｏ＋Ｃｒ≧０．２％、かつＣｒ≦０．５％、Ｍｏ≦０．５％、を満たす範囲でＣ、Ｓ、Ｎ、Ｔｉを含有し残部がＦｅ及び不可避的不純物からなる鋼であって、そのミクロ組織が、フェライト、またはフェライトおよびベイナイトからなるバーリング性高強度鋼板がバーリング性は非常に優れるものの溶接熱影響部が著しく軟化することを知見した。さらに上記バーリング性高強度鋼板の溶接熱影響部軟化の原因が溶接温度履歴によるミクロ組織の焼き戻しによるものであることを突き止め、耐軟化性を向上させるためにはＣｒ、Ｍｏの複合添加が非常に有効であることを新たに見出し、本発明をなしたものである。 The present inventors improve the softening resistance of the welding heat affected zone of a burring high strength steel sheet with a mind on the manufacturing process of a thin steel sheet produced on an industrial scale by manufacturing equipment that is currently commonly used. I did my utmost research. As a result, C: 0.01 to 0.1%, Si: 0.01 to 2%, Mn: 0.05 to 3%, P ≦ 0.1%, S ≦ 0.03%, Al: 0. 005-1%, N: 0.0005-0.005%, Ti: 0.05-0.5%, and 0 <C- (12 / 48Ti-12 / 14N-12 / 32S) ≤0. 0.05%, Mo + Cr ≧ 0.2%, and Cr ≦ 0.5%, Mo ≦ 0.5%, in the range that satisfies the following conditions: C, S, N, Ti, with the balance being Fe and unavoidable impurities. It has been found that a high-strength burring steel sheet composed of ferrite or ferrite and bainite has a very good burring property, but has a significantly softened weld heat affected zone. Furthermore, it was found that the cause of the softening of the weld heat affected zone of the burring high strength steel sheet was due to the tempering of the microstructure due to the welding temperature history, and in order to improve the softening resistance, the addition of Cr and Mo was very important. It is newly found that the present invention is effective, and the present invention has been made.
即ち、本発明の要旨は、以下の通りである。 That is, the gist of the present invention is as follows.
（１） 質量％にて、Ｃ：０．０１〜０．１％、Ｓｉ：０．０１〜２％、Ｍｎ：０．０５〜３％、Ｐ≦０．１％、Ｓ≦０．０３％、Ａｌ：０．００５〜１％、Ｎ：０．０００５〜０．００５％、Ｔｉ：０．０５〜０．５％、を含み、さらに０％＜Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％、さらに、Ｍｏ＋Ｃｒ≧０．２％、かつＣｒ≦０．５％、Ｍｏ≦０．５％、を満たす範囲でＣ、Ｓ、Ｎ、Ｔｉ、Ｃｒ、Ｍｏを含有し残部がＦｅ及び不可避的不純物からなる鋼であって、そのミクロ組織が、フェライト、またはフェライトおよびベイナイトからなることを特徴とする溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (1) In mass%, C: 0.01 to 0.1%, Si: 0.01 to 2%, Mn: 0.05 to 3%, P ≦ 0.1%, S ≦ 0.03% , Al: 0.005 to 1%, N: 0.0005 to 0.005%, Ti: 0.05 to 0.5%, and further 0% <C− (12 / 48Ti-12 / 14N−). 12 / 32S) ≦ 0.05%, and furthermore, C, S, N, Ti, Cr and Mo within a range satisfying Mo + Cr ≧ 0.2%, Cr ≦ 0.5%, and Mo ≦ 0.5%. A burring high-strength steel sheet excellent in softening resistance of a weld heat-affected zone, characterized in that the steel contains Fe and unavoidable impurities, and the microstructure of the steel is ferrite or ferrite and bainite. .
（２） 前記鋼がさらに、質量％にて、Ｎｂ：０．０１〜０．５％を含み、さらに０＜Ｃ−（１２／４８Ｔｉ＋１２／９３Ｎｂ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％、を満たす範囲でＮｂを含有し残部がＦｅ及び不可避的不純物からなる鋼であることを特徴とする溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (2) The steel further contains Nb: 0.01 to 0.5% by mass%, and 0 <C− (12 / 48Ti + 12 / 93Nb-12 / 14N-12 / 32S) ≦ 0.05. %. A burring high-strength steel sheet excellent in softening resistance of a heat-affected zone of a weld, wherein the steel contains Nb in a range satisfying% and the balance is Fe and inevitable impurities.
（３） （１）又は（２）に記載の鋼が、さらに、質量％にて、Ｃａ：０．０００５〜０．００２％、ＲＥＭ：０．０００５〜０．０２％の一種または二種を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (3) The steel according to (1) or (2) further contains one or two types of Ca: 0.0005 to 0.002% and REM: 0.0005 to 0.02% by mass%. A burring high-strength steel sheet excellent in softening resistance of a heat affected zone of welding, characterized by containing.
（４） （１）ないし（３）のいずれか１項に記載の鋼が、さらに、質量％にて、Ｃｕ：０．２〜１．２％を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (4) The welding heat effect, wherein the steel according to any one of (1) to (3) further contains Cu: 0.2 to 1.2% by mass%. Burring high strength steel sheet with excellent softening resistance in the part.
（５） （１）ないし（４）のいずれか１項に記載の鋼が、さらに、質量％にて、Ｎｉ：０．１〜０．６％を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (5) The effect of welding heat, wherein the steel according to any one of (1) to (4) further contains, by mass%, Ni: 0.1 to 0.6%. Burring high strength steel sheet with excellent softening resistance in the part.
（６） （１）ないし（５）のいずれか１項に記載の鋼が、さらに、質量％にて、Ｂ：０．０００２〜０．００２％を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (6) The welding heat effect, wherein the steel according to any one of (1) to (5) further contains, by mass%, B: 0.0002 to 0.002%. Burring high strength steel sheet with excellent softening resistance in the part.
（７） （１）ないし（６）のいずれか１項に記載の自動車用薄鋼板に亜鉛めっきが施されていることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 (7) A high burring property excellent in softening resistance of the weld heat affected zone, characterized in that the thin steel sheet for automobiles according to any one of (1) to (6) is galvanized. Strength steel plate.
（８） （１）ないし（６）のいずれか１項に記載の薄鋼板を得るために該成分を有する鋼片の熱間圧延に際して仕上圧延をＡｒ3変態点温度＋３０℃以上の温度域で終了し、その後１０秒以内に冷却終了までの平均冷却速度が５０℃／秒以上の冷却速度で７００℃以下の温度域まで冷却し、３５０℃以上６５０℃以下の巻き取り温度にて巻き取ることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板の製造方法。 (8) In order to obtain the thin steel sheet according to any one of (1) to (6), the finish rolling is performed in the temperature range of the Ar 3 transformation point temperature + 30 ° C. or more at the time of hot rolling of a slab having the component. After finishing, cooling within 10 seconds at an average cooling rate of 50 ° C./sec or more to a temperature range of 700 ° C. or less, and winding at a winding temperature of 350 ° C. to 650 ° C. A method for producing a burring high-strength steel sheet having excellent softening resistance in a heat affected zone by welding.
（９） （１）ないし（６）のいずれか１項に記載の薄鋼板を得るために該成分を有する鋼片を熱間圧延、酸洗、冷間圧延後、８００℃以上の温度域で５〜１５０秒間保持し、その後平均冷却速度が５０℃／秒以上の冷却速度で７００℃以下の温度域まで冷却する工程の熱処理をすることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板の製造方法。 (9) In order to obtain the thin steel sheet according to any one of (1) to (6), a slab having the component is hot-rolled, pickled, and cold-rolled, and then subjected to a temperature range of 800 ° C or higher. Holding for 5 to 150 seconds, and then performing a heat treatment in a step of cooling to a temperature range of 700 ° C. or less at a cooling rate of 50 ° C./sec or more to the softening resistance of the heat affected zone. Excellent burring high strength steel sheet manufacturing method.
（１０） （８）に記載の製造方法に際し、熱間圧延工程終了後に亜鉛めっき浴中に浸積させて鋼板表面を亜鉛めっきすることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板の製造方法。 (10) In the production method described in (8), the steel sheet surface is galvanized by dipping in a galvanizing bath after completion of the hot rolling step, and the welding heat affected zone has excellent softening resistance. For producing high-strength burring steel sheets.
（１１） （９）に記載の製造方法に際し、熱処理工程終了後、亜鉛めっき浴中に浸積させて鋼板表面を亜鉛めっきすることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板の製造方法。 (11) In the production method according to (9), after the heat treatment step, the steel sheet is immersed in a galvanizing bath to galvanize the surface of the steel sheet. A method for producing a high strength burring steel sheet.
（１２） （１０）又は（１１）に記載の製造方法に際し、亜鉛めっき浴中に浸積して亜鉛めっき後、合金化処理することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板の製造方法。 (12) The method of manufacturing according to (10) or (11), characterized in that it is immersed in a galvanizing bath, galvanized, and then subjected to an alloying treatment. For producing high-strength burring steel sheets.
以上詳述したように、本発明は、溶接熱影響部の耐軟化性に優れた引張強度５４０ＭＰａ以上のバーリング性高強度鋼板およびその製造方法に関するものであり、これらの薄鋼板を用いることにより、成形後にスポット、アーク、プラズマ、レーザー等により溶接される場合や、これら溶接後に成形される場合において溶接熱影響部の耐軟化性の大幅な改善が期待できるため、本発明は、工業的価値が高い発明であると言える。 As described in detail above, the present invention relates to a high-strength burring steel sheet having a tensile strength of 540 MPa or more and excellent in softening resistance of a weld heat-affected zone, and a method for manufacturing the same. By using these thin steel sheets, Spots, arcs, plasma after forming, when welded by laser, etc., or when formed after these weldings, a significant improvement in the softening resistance of the heat affected zone can be expected. It can be said that this is a high invention.
まず、溶接熱影響部の耐軟化性に及ぼすＣ*量（Ｃ*＝Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）：以下Ｃ*と標記する。）およびＣｒ、Ｍｏ含有量の影響についての調査を行った。そのための供試材は、次のようにして準備した。すなわち、０．０５％Ｃ−１．０％Ｓｉ−１．４％Ｍｎ−０．０１％Ｐ−０.００１％ＳをベースにＣ*量（Ｔｉ、Ｎ含有量）およびＣｒ＋Ｍｏ量を変化させて成分調整し溶製した鋳片を熱間圧延して常温で巻き取り、５５０℃で１時間等温保持した後、炉冷する熱処理を施した。これらの鋼板についてアーク溶接部硬度測定を行った結果を図２に示す。 First, the amount of C * (C * = C- (12 / 48Ti-12 / 14N-12 / 32S): hereinafter referred to as C * ) that affects the softening resistance of the heat-affected zone of welding and the contents of Cr and Mo are described below. The impact was investigated. Test materials for that purpose were prepared as follows. That is, based on 0.05% C-1.0% Si-1.4% Mn-0.01% P-0.001% S, the C * content (Ti and N content) and the Cr + Mo content were changed. The slab thus prepared and melted was hot-rolled, wound at room temperature, kept at 550 ° C. for 1 hour, and then subjected to a heat treatment for furnace cooling. FIG. 2 shows the results of arc hardness measurement of these steel sheets.
ここで、この結果より、Ｃ*量およびＣｒ＋Ｍｏ量と溶接熱影響部の軟化程度ΔＨｖ（ΔＨｖ＝Ｈｖ（母材硬度平均値）−Ｈｖ（溶接熱影響最軟化部硬度）と定義する：図１参照）には強い相関があり、Ｃ*量が０より大きく０．０５％以下かつＣｒ＋Ｍｏ量が０．２％以上で溶接熱影響部の軟化が著しく抑制されることを新規に知見した。 Here, based on the results, the amounts of C * and Cr + Mo and the degree of softening of the weld heat affected zone are defined as ΔHv (ΔHv = Hv (average base material hardness) −Hv (weld heat affected softest part hardness): FIG. Has a strong correlation, and newly found that softening of the heat affected zone is significantly suppressed when the amount of C * is greater than 0 and equal to or less than 0.05% and the amount of Cr + Mo is equal to or greater than 0.2%.
このメカニズムは必ずしも明らかではないが、ベイニティックなミクロ組織により強度を得ている材料は、アーク溶接等の溶接熱サイクルでその熱影響部が軟化する場合がある。ＭｏもしくはＣｒは溶接のような短時間の熱サイクルでも、Ｃ等の元素とクラスタリングもしくは析出して強度を上昇させ、結果として熱影響部の軟化を抑制したと推測される。ただし、ＭｏとＣｒの含有量の合計が０．２％未満ではこの効果が失われる。 Although the mechanism is not always clear, the heat-affected zone of a material having a strength obtained by a bainitic microstructure may be softened by a welding heat cycle such as arc welding. It is presumed that Mo or Cr increases the strength by clustering or precipitating with elements such as C even in a short heat cycle such as welding, thereby suppressing the softening of the heat-affected zone. However, if the total content of Mo and Cr is less than 0.2%, this effect is lost.
一方、ＭｏもしくはＣｒ炭化物等を得るためには、ＴｉＣ等の高温で析出する炭化物で固定される当量以上のＣを含有しなければならない。従って、Ｃ＊≦０でこの効果は失われる。 On the other hand, in order to obtain Mo or Cr carbide or the like, it is necessary to contain an equivalent or more of C fixed by carbide precipitated at a high temperature such as TiC. Therefore, this effect is lost when C * ≦ 0.
なお、アーク溶接の溶接熱影響部の硬度測定はについては、ＪＩＳ Ｚ ３１０１記載の１号試験片にて、ＪＩＳ Ｚ ２２４４記載の試験方法に準じてで測定した。ただし、アーク溶接は、シールドガス：ＣＯ2、ワイヤ：日鐵溶接工業（株）製ＹＭ−６０Ｃφ１．２ｍｍを用い、溶接速度：１００ｃｍ／分、溶接電流：２６０±１０Ａ、溶接電圧：２６±１Ｖ、供試材の板厚は２．６ｍｍとし、硬度測定位置は、表面より０．２５ｍｍ、測定間隔は、０．５ｍｍで、試験力は９８Ｎとした。 In addition, about the hardness measurement of the welding heat affected zone of arc welding, it measured according to the test method of JISZ2244 using the 1st test piece of JISZ3101. However, in arc welding, shield gas: CO 2 , wire: YM-60C φ1.2 mm manufactured by Nippon Steel Welding Co., Ltd., welding speed: 100 cm / min, welding current: 260 ± 10 A, welding voltage: 26 ± 1 V The thickness of the test material was 2.6 mm, the hardness measurement position was 0.25 mm from the surface, the measurement interval was 0.5 mm, and the test force was 98 N.
次に、本発明における鋼板のミクロ組織について説明する。 Next, the microstructure of the steel sheet according to the present invention will be described.
鋼板のミクロ組織は、優れたバーリング加工性を確保するためにフェライト単相が望ましい。ただし、必要に応じ一部ベイナイトを含むことを許容するものであるが、良好なバーリング加工性を確保するためには、ベイナイトの体積分率は１０％以下が望ましい。なお、ここで言うフェライトとはベイニティックフェライトおよびアシュキュラーフェライト組織も含む。また、ベイナイトとは透過型電子顕微鏡にて薄膜を観察した場合フェライトラス間にセメンタイト等の炭化物を含むかもしくはフェライトラス内にセメンタイト等の炭化物を含む組織である。一方、ベイニティックフェライトおよびアシュキュラーフェライト組織とはＴｉ、Ｎｂの炭窒化物以外はフェライトラス内およびフェライトラス間に炭化物を含まない組織と定義する。 The microstructure of the steel sheet is desirably a ferrite single phase in order to ensure excellent burring workability. However, although it is allowed to partially contain bainite as needed, the volume fraction of bainite is desirably 10% or less in order to ensure good burring workability. The ferrite referred to here includes bainitic ferrite and ashular ferrite structures. Further, bainite is a structure containing carbide such as cementite between ferrite laths or containing carbide such as cementite in ferrite laths when a thin film is observed with a transmission electron microscope. On the other hand, bainitic ferrite and ashular ferrite microstructures are defined as microstructures that do not contain carbides in ferrite laths and between ferrite laths except for carbonitrides of Ti and Nb.
また、不可避的なマルテンサイト、残留オーステナイトおよびパーライトを含むことを許容するものであるが、良好なバーリング性を確保するためには、残留オーステナイトおよびマルテンサイトを合わせた体積分率は５％未満が望ましい。さらに、良好な疲労特性を確保するためには、粗大な炭化物を含むパーライトの体積分率は５％以下が望ましい。また、ここで、フェライト、ベイナイト、残留オーステナイト、パーライト、マルテンサイトの体積分率とは鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より切出した試料を圧延方向断面に研磨し、ナイタール試薬を用いてエッチングし、光学顕微鏡を用い２００〜５００倍の倍率で観察された板厚の１／４ｔにおけるミクロ組織の面積分率で定義される。 In addition, it is allowed to contain unavoidable martensite, residual austenite and pearlite. However, in order to secure good burring properties, the combined volume fraction of residual austenite and martensite is less than 5%. desirable. Further, in order to secure good fatigue properties, the volume fraction of pearlite containing coarse carbides is desirably 5% or less. Here, the volume fractions of ferrite, bainite, retained austenite, pearlite, and martensite are as follows. A sample cut from a 1 / 4W or 3 / 4W position of the steel sheet width is polished into a section in the rolling direction, and a nital reagent is used. And defined by the area fraction of the microstructure at 1 / 4t of the plate thickness observed at a magnification of 200 to 500 times using an optical microscope.
次に、本発明の化学成分の限定理由について説明する。 Next, the reasons for limiting the chemical components of the present invention will be described.
Ｃは、本発明における最も重要な元素の一つである。すなわち、Ｃは、溶接のような短時間の熱サイクルでもＭｏもしくはＣｒとクラスタリングもしくは析出して溶接熱影響部の軟化を抑制する効果がある。ただし、０．１％超含有していると加工性及び溶接性が劣化するので、０．１％以下とする。また０．０１％未満であると強度が低下するので０．０１％以上とする。 C is one of the most important elements in the present invention. That is, C has the effect of suppressing or softening the weld heat affected zone by clustering or precipitating with Mo or Cr even in a short heat cycle such as welding. However, if the content exceeds 0.1%, workability and weldability deteriorate, so the content is set to 0.1% or less. If it is less than 0.01%, the strength is reduced.
Ｓｉは、固溶強化元素として強度上昇に有効である。所望の強度を得るためには、０．０１％以上含有する必要がある。しかし、２％超含有すると加工性が劣化する。そこで、Ｓｉの含有量は０．０１％以上、２％以下とする。 Si is effective for increasing the strength as a solid solution strengthening element. In order to obtain a desired strength, it is necessary to contain 0.01% or more. However, if the content exceeds 2%, the workability deteriorates. Therefore, the content of Si is set to 0.01% or more and 2% or less.
Ｍｎは、固溶強化元素として強度上昇に有効である。所望の強度を得るためには、０．０５％以上必要である。また、Ｍｎ以外にＳによる熱間割れの発生を抑制するＴｉなどの元素が十分に添加されない場合には質量％でＭｎ／Ｓ≧２０となるＭｎ量を添加することが望ましい。一方、３％超添加するとスラブ割れを生ずるため、３％以下とする。 Mn is effective for increasing strength as a solid solution strengthening element. To obtain a desired strength, 0.05% or more is required. When an element such as Ti that suppresses hot cracking due to S other than Mn is not sufficiently added, it is preferable to add an Mn amount that satisfies Mn / S ≧ 20 in mass%. On the other hand, if added over 3%, slab cracks occur, so the content is set to 3% or less.
Ｐは、不純物であり低いほど望ましく、０．１％超含有すると加工性や溶接性に悪影響を及ぼすとともに疲労特性も低下させるので、０．１％以下とする。
Ｓは、多すぎると熱間圧延時の割れを引き起こすので極力低減させるべきであるが、０．０３％以下ならば許容できる範囲である。
Ａｌは、溶鋼脱酸のために０．００５％以上添加する必要があるが、コストの上昇を招くため、その上限を１％とする。また、あまり多量に添加すると、非金属介在物を増大させ伸びを劣化させるので望ましくは０．５％以下とする。
P is an impurity and is preferably as low as possible, and if it exceeds 0.1%, it adversely affects workability and weldability and also deteriorates fatigue characteristics.
If S is too large, it causes cracking during hot rolling, so it should be reduced as much as possible, but if it is 0.03% or less, it is in an acceptable range.
Al needs to be added in an amount of 0.005% or more for deoxidation of molten steel. However, since the cost is increased, the upper limit is set to 1%. Further, if added in an excessively large amount, nonmetallic inclusions increase and elongation deteriorates. Therefore, the content is desirably 0.5% or less.
Ｎは、Ｃよりも高温にてＴｉおよびＮｂと析出物を形成し、所望のＣを固定するのに有効なＴｉおよびＮｂを減少させる。従って極力低減させるべきであるが、０．００５％以下ならば許容できる範囲である。 N forms a precipitate with Ti and Nb at a higher temperature than C, and reduces Ti and Nb which are effective for fixing desired C. Therefore, it should be reduced as much as possible, but if it is 0.005% or less, it is within an acceptable range.
Ｔｉは、本発明における最も重要な元素の一つである。すなわち、Ｔｉは析出強化により鋼板の強度上昇に寄与する。ただし、０．０５％未満ではこの効果が不十分であり、０．５％超含有してもその効果が飽和するだけでなく合金コストの上昇を招く。従ってＴｉの含有量は０．０５％以上、０．５％以下とする。さらに、バーリング加工性を劣化させるセメンタイト等の炭化物の原因となるＣを析出固定し、バーリング加工性の向上に寄与するためには、Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％の条件を満たすことが必要である。一方、溶接熱影響部の軟化抑制の面からは、ＭｏもしくはＣｒをクラスタリングもしくは析出させるに十分な固溶Ｃが必要であるので、０＜Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）とする。 Ti is one of the most important elements in the present invention. That is, Ti contributes to an increase in the strength of the steel sheet by precipitation strengthening. However, if the content is less than 0.05%, this effect is insufficient. If the content exceeds 0.5%, not only the effect is saturated, but also the alloy cost is increased. Therefore, the content of Ti is set to 0.05% or more and 0.5% or less. Further, in order to precipitate and fix C, which causes carbide such as cementite, which deteriorates burring workability, and to contribute to improvement in burring workability, C- (12 / 48Ti-12 / 14N-12 / 32S) ≦ It is necessary to satisfy the condition of 0.05%. On the other hand, from the aspect of suppressing the softening of the weld heat affected zone, a solid solution C sufficient for clustering or precipitating Mo or Cr is necessary, so that 0 <C- (12 / 48Ti-12 / 14N-12 / 32S ).
Ｍｏ、Ｃｒは、本発明の最も重要な元素の一つであり、溶接のような短時間の熱サイクルでも、Ｃ等の元素とクラスタリングもしくは析出して熱影響部の軟化を抑制する。ただし、ＭｏとＣｒの含有量の合計が０．２％未満ではこの効果が失われる。また、それぞれ、０．５％超含有してもその効果が飽和するので、それぞれ、Ｍｏ≦０．５％、Ｃｒ≦０．５％とする。 Mo and Cr are one of the most important elements of the present invention, and suppress the softening of the heat affected zone by clustering or precipitating with elements such as C even in a short heat cycle such as welding. However, if the total content of Mo and Cr is less than 0.2%, this effect is lost. The effect is saturated even if the content exceeds 0.5%, so that Mo ≦ 0.5% and Cr ≦ 0.5%, respectively.
Ｎｂは、Ｔｉ同様に析出強化により鋼板の強度上昇に寄与する。ただし、０．０１％未満ではこの効果が不十分であり、０．５％超含有してもその効果が飽和するだけでなく合金コストの上昇を招く。従ってＮｂの含有量は０．０１％以上、０．５％以下とする。さらに、バーリング加工性を劣化させるセメンタイト等の炭化物の原因となるＣを析出固定し、Ｃ−（１２／４８Ｔｉ＋１２／９３Ｎｂ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％の条件を満たすことが必要である。一方、溶接熱影響部の軟化抑制の面からは、ＭｏもしくはＣｒをクラスタリングもしくは析出させるに十分な固溶Ｃが必要であるので、０＜Ｃ−（１２／４８Ｔｉ＋１２／９３Ｎｂ−１２／１４Ｎ−１２／３２Ｓ）とする。 Nb contributes to an increase in the strength of the steel sheet by precipitation strengthening like Ti. However, if the content is less than 0.01%, this effect is insufficient. If the content exceeds 0.5%, not only the effect is saturated, but also the alloy cost is increased. Therefore, the content of Nb is set to 0.01% or more and 0.5% or less. Further, C which causes carbide such as cementite which deteriorates burring workability is deposited and fixed, and the condition of C- (12 / 48Ti + 12 / 93Nb-12 / 14N-12 / 32S) ≦ 0.05% is satisfied. is necessary. On the other hand, from the viewpoint of suppressing the softening of the weld heat affected zone, a solid solution C sufficient for clustering or precipitating Mo or Cr is necessary, so that 0 <C− (12 / 48Ti + 12 / 93Nb−12 / 14N−12). / 32S).
ＣａおよびＲＥＭは、破壊の起点となったり、加工性を劣化させる非金属介在物の形態を変化させて無害化する元素である。ただし、０．００５％未満添加してもその効果がなく、Ｃａならば０．０２％超、ＲＥＭならば０．２％超添加してもその効果が飽和するのでＣａ＝０．００５〜０．０２％、ＲＥＭ＝０．００５〜０．２％添加することが好ましい。 Ca and REM are elements that become the starting point of destruction and change the form of nonmetallic inclusions that degrade workability and render them harmless. However, if less than 0.005% is added, there is no effect. If Ca is added more than 0.02%, and if REM is added more than 0.2%, the effect is saturated. 0.02%, and REM = 0.005 to 0.2%.
Ｃｕは、固溶状態で疲労特性を改善する効果がある。ただし、０．２％未満では、その効果は少なく、１．２％を超えて含有すると巻取り中に析出して析出強化により鋼板の静的強度が著しく上昇するため、加工性が著しく劣化することになる。また、このようなＣｕの析出強化では、疲労限は静的強度の上昇ほどには向上しないので疲労限度比が低下してしまう。そこで、Ｃｕの含有量は０．２〜１．２％の範囲とする。 Cu has an effect of improving fatigue characteristics in a solid solution state. However, if the content is less than 0.2%, the effect is small. If the content exceeds 1.2%, precipitation occurs during winding and the precipitation strengthening significantly increases the static strength of the steel sheet. Will be. In addition, with such Cu precipitation strengthening, the fatigue limit does not improve as much as the increase in static strength, so that the fatigue limit ratio decreases. Therefore, the content of Cu is set in the range of 0.2 to 1.2%.
Ｎｉは、Ｃｕ含有による熱間脆性防止のために必要に応じ添加する。ただし、０．１％未満ではその効果が少なく、１％を超えて添加してもその効果が飽和するので、０．１〜１％とする。 Ni is added as necessary to prevent hot brittleness due to the inclusion of Cu. However, if the content is less than 0.1%, the effect is small, and if the content exceeds 1%, the effect is saturated. Therefore, the content is set to 0.1 to 1%.
Ｂは、固溶Ｃ量の減少が原因と考えられるＰによる粒界脆化を抑制することによって疲労限を上昇させる効果があるので必要に応じ添加する。さらに、母材強度が６４０ＭＰａ以上である場合、溶接熱影響部のうちα→γ→α変態が起こる熱履歴を受ける部位において低Ｃｅｑ故に焼が入らず、軟化する恐れがある場合に焼入れ性を向上させるＢを添加することにより、当該部位での軟化を抑制し、継手の破断形態を溶接部から、母材部へ遷移させる効果があるので必要に応じて添加する。ただし、０．０００２％未満ではその効果を得るために不十分であり、０．００２％超添加するとスラブ割れが起こる。よって、Ｂの添加は、０．０００２％以上、０．００２％以下とする。 B is added as necessary because it has the effect of increasing the fatigue limit by suppressing grain boundary embrittlement due to P, which is considered to be caused by the decrease in the amount of solute C. Further, when the base metal strength is 640 MPa or more, hardening does not occur due to low Ceq in a part of the heat affected zone of the weld which undergoes a heat history in which α → γ → α transformation occurs, and the hardenability is increased when there is a possibility of softening. Addition of B to be improved has the effect of suppressing softening at the site and transitioning the fracture mode of the joint from the welded portion to the base material portion, so that it is added as necessary. However, if it is less than 0.0002%, it is insufficient to obtain the effect, and if it exceeds 0.002%, slab cracking occurs. Therefore, the addition of B is set to 0.0002% or more and 0.002% or less.
さらに、強度を付与するために、Ｖ、Ｚｒの析出強化もしくは固溶強化元素の一種または二種以上を添加してもよい。ただし、それぞれ、０．０２％、０．０２％未満ではその効果を得ることができない。また、それぞれ、０．２％、０．２％を超え添加してもその効果は飽和する。 Further, in order to impart strength, one or two or more elements of precipitation strengthening or solid solution strengthening of V and Zr may be added. However, the effects cannot be obtained if the content is less than 0.02% and 0.02%, respectively. The effect is saturated even if they are added in excess of 0.2% and 0.2%, respectively.
なお、これらを主成分とする鋼にＳｎ、Ｃｏ、Ｚｎ、Ｗ、Ｍｇを合計で１％以下含有しても構わない。しかしながらＳｎは熱間圧延時に疵が発生する恐れがあるので０．０５％以下が望ましい。 In addition, steel containing these as main components may contain Sn, Co, Zn, W, and Mg in a total of 1% or less. However, Sn may be flawed at the time of hot rolling, so that 0.05% or less is desirable.
次に、本発明の製造方法の限定理由について、以下に詳細に述べる。 Next, the reasons for limiting the production method of the present invention will be described in detail below.
本発明は、鋳造後、熱間圧延後冷却ままもしくは熱間圧延後、熱間圧延後冷却・酸洗し冷延した後に熱処理、あるいは熱延鋼板もしくは冷延鋼板を溶融めっきラインにて熱処理を施したまま、更にはこれらの鋼板に別途表面処理を施すことによっても得られる。 The present invention, after casting, hot rolling after cold rolling or hot rolling, heat treatment after cooling, pickling and cold rolling after hot rolling, or heat treatment of hot-rolled steel sheet or cold-rolled steel sheet in a hot dip coating line As it is, it can also be obtained by subjecting these steel sheets to a separate surface treatment.
本発明において熱間圧延に先行する製造方法は特に限定するものではない。すなわち、高炉や電炉等による溶製に引き続き各種の２次製錬で目的の成分含有量になるように成分調整を行い、次いで通常の連続鋳造、インゴット法による鋳造の他、薄スラブ鋳造などの方法で鋳造すればよい。原料にはスクラップを使用しても構わない。連続鋳造よって得たスラブの場合には高温鋳片のまま熱間圧延機に直送してもよいし、室温まで冷却後に加熱炉にて再加熱した後に熱間圧延してもよい。 In the present invention, the production method prior to hot rolling is not particularly limited. In other words, following smelting in a blast furnace or electric furnace, the components are adjusted in the various secondary refining processes so that the desired component content is obtained, and then normal continuous casting, casting by ingot method, thin slab casting, etc. It may be cast by a method. Scrap may be used as a raw material. In the case of a slab obtained by continuous casting, the slab may be directly sent to a hot rolling mill as it is, or may be cooled to room temperature and then re-heated in a heating furnace before hot rolling.
再加熱温度については特に制限はないが、１４００℃以上であると、スケールオフ量が多量になり歩留まりが低下するので、再加熱温度は１４００℃未満が望ましい。また、１０００℃未満の加熱はスケジュール上操業効率を著しく損なうため、再加熱温度は１０００℃以上が望ましい。さらには、１１００℃未満での加熱はＴｉおよび／またはＮｂを含む析出物がスラブ中で再溶解せず粗大化し析出強化能を失うばかりでなくバーリング加工性にとって望ましいサイズと分布のＴｉおよび／またはＮｂを含む析出物が析出しなくなるので、再加熱温度は１１００℃以上が望ましい。 The reheating temperature is not particularly limited, but if it is 1400 ° C. or more, the scale-off amount becomes large and the yield decreases, so the reheating temperature is desirably less than 1400 ° C. Further, since the heating at a temperature lower than 1000 ° C. significantly impairs the operation efficiency on a schedule, the reheating temperature is preferably 1000 ° C. or higher. Furthermore, heating below 1100 ° C. not only causes the precipitates containing Ti and / or Nb not to be redissolved in the slab and coarsens to lose the precipitation strengthening ability, but also to have the desired size and distribution of Ti and / or distribution for burring workability. Since the precipitate containing Nb is not deposited, the reheating temperature is desirably 1100 ° C. or higher.
熱間圧延工程は、粗圧延を終了後、仕上げ圧延を行うが、粗圧延後または、それに続くデスケーリング後にシートバーを接合し、連続的に仕上げ圧延をしてもよい。その際に粗バーを一旦コイル状に巻き、必要に応じて保温機能を有するカバーに格納し、再度巻き戻してから接合を行ってもよい。また、その後の仕上げ圧延はデスケーリング後に再びスケールが生成してしまうのを防ぐために５秒以内に行うのが望ましい。 In the hot rolling step, the finish rolling is performed after the rough rolling is completed, but the sheet bar may be joined after the rough rolling or after the subsequent descaling, and the finish rolling may be continuously performed. At that time, the coarse bar may be temporarily wound in a coil shape, stored in a cover having a heat retaining function as necessary, and then re-wound before joining. Further, the subsequent finish rolling is desirably performed within 5 seconds in order to prevent scale from being formed again after descaling.
仕上げ圧延は、最終パス温度（ＦＴ）がＡｒ3変態点＋３０℃以上の温度域で終了する必要がある。これは、熱間圧延後の冷却工程においてバーリング加工性にとって好ましいベイニティックなフェライト、またはフェライトおよびベイナイトを得るためにγ→α変態が低温で起こることが必要であるが、最終パス温度（ＦＴ）がＡｒ3変態点＋３０℃未満の温度域ではひずみ誘起によるフェライト変態核生成が起こり、ポリゴナルで粗大なフェライトが生成してしまう懸念がある。仕上げ温度の上限は本発明の効果を得るためには特に定める必要はないが、操業上スケール疵が発生する可能性があるため、１１００℃以下とすることが好ましい。ここでＡｒ3変態点温度とは、例えば以下の計算式により鋼成分との関係で簡易的に示される。すなわち
Ａｒ3＝９１０−３１０×％Ｃ＋２５×％Ｓｉ−８０×％Ｍｎ
Finish rolling must be completed in a temperature range where the final pass temperature (FT) is equal to or higher than the Ar 3 transformation point + 30 ° C. This is because the γ → α transformation needs to occur at low temperature to obtain bainitic ferrite, or ferrite and bainite, which is favorable for burring workability in the cooling step after hot rolling, but the final pass temperature (FT ) In a temperature range lower than the Ar 3 transformation point + 30 ° C., there is a concern that ferrite transformation nucleation occurs due to strain induction and polygonal and coarse ferrite is produced. The upper limit of the finishing temperature is not particularly required to obtain the effect of the present invention, but is preferably 1100 ° C. or lower because scale flaws may occur during operation. Here, the Ar 3 transformation point temperature is simply indicated in relation to the steel composition by the following calculation formula, for example. That Ar 3 = 910-310 ×% C + 25 ×% Si-80 ×% Mn
仕上圧延を終了した後は、指定の巻取温度（ＣＴ）まで冷却するが、その冷却開始までの時間は１０秒以内とする。これは冷却開始までの時間が１０秒超であると圧延直後に再結晶したオーステナイト粒が粗大化してしまいγ→α変態後のフェライト粒が粗大化してしまう懸念があるからである。次に冷却終了までの平均冷却速度であるが、５０℃／秒以上が必要である。これは冷却終了までの平均冷却速度が5０℃／秒未満であるとバーリング加工性にとって好ましいベイニティックなフェライト、またはフェライトおよびベイナイトの体積分率が減少してします恐れがあるからである。また、冷却速度の上限は実際の工場設備能力等を考慮すると５００℃／秒以下である。冷却終了温度は７００℃以下の温度域であることが必要である。これは冷却終了温度が７００℃超であるとバーリング加工性にとって好ましいベイニティックなフェライト、またはフェライトおよびベイナイト以外のミクロ組織が生成してしまう怖れがあるからである。冷却終了温度の下限は本発明の効果を得るためには特に定める必要はない。ただし、巻き取り温度以下には本発明のプロセス上ありえない。冷却終了後から巻き取りまでのの工程については特に定めないが、必要に応じて巻き取り温度まで冷却してもよいが、この場合、熱ひずみによる板そりが懸念されることから、３００℃／ｓ以下とすることが望ましい。 After finishing the finish rolling, it is cooled to a specified winding temperature (CT), and the time until the start of cooling is within 10 seconds. This is because if the time until the start of cooling exceeds 10 seconds, austenite grains recrystallized immediately after rolling become coarse, and there is a concern that ferrite grains after the γ → α transformation become coarse. Next, the average cooling rate up to the end of cooling is required to be 50 ° C./sec or more. This is because if the average cooling rate until the end of cooling is less than 50 ° C./sec, the volume fraction of bainitic ferrite or ferrite and bainite, which is preferable for burring workability, may decrease. Further, the upper limit of the cooling rate is 500 ° C./sec or less in consideration of the actual plant facility capacity and the like. The cooling end temperature needs to be in a temperature range of 700 ° C. or less. This is because if the cooling end temperature is higher than 700 ° C., there is a fear that bainitic ferrite or a microstructure other than ferrite and bainite which is preferable for burring workability may be formed. The lower limit of the cooling end temperature does not need to be particularly determined in order to obtain the effects of the present invention. However, below the winding temperature, it is impossible in the process of the present invention. The process from the end of cooling to the winding is not particularly limited, but may be cooled to the winding temperature if necessary. s or less is desirable.
次に巻取温度が３５０℃未満では十分なＴｉおよび／またはＮｂを含む析出物が生じなくなり、強度低下が懸念される、６５０℃超ではＴｉおよび／またはＮｂを含む析出物のサイズが粗大化し析出強化による強度上昇に寄与しなくなるばかりでなく、析出物が大きすぎると析出物と母相の界面にボイドが生じやすくなり、穴拡性が低下する恐れがある。従って巻取温度は３５０℃〜６５０℃とする。さらに、巻取り後の冷却速度は特に限定しないが、Ｃｕを１％以上添加した場合、巻取温度（ＣＴ）が４５０℃超であると巻取り後にＣｕが析出して加工性が劣化するばかりでなく、疲労特性向上に有効な固溶状態のＣｕが失われる恐れがあるので、巻取温度（ＣＴ）が４５０℃超の場合、巻取り後の冷却速度は２００℃までを３０℃／ｓ以上とすることが望ましい。 Next, if the winding temperature is lower than 350 ° C., a precipitate containing sufficient Ti and / or Nb is not generated, and there is a concern that the strength is reduced. If it exceeds 650 ° C., the size of the precipitate containing Ti and / or Nb becomes coarse. Not only does it not contribute to the increase in strength due to precipitation strengthening, but if the precipitate is too large, voids are likely to be formed at the interface between the precipitate and the mother phase, and the hole expandability may decrease. Therefore, the winding temperature is set to 350 ° C to 650 ° C. Further, the cooling rate after winding is not particularly limited, but when Cu is added at 1% or more, if the winding temperature (CT) is more than 450 ° C, Cu precipitates after winding and the workability is deteriorated. In addition, when the winding temperature (CT) is higher than 450 ° C., the cooling rate after winding is reduced from 200 ° C. to 30 ° C./s, since Cu in the solid solution effective for improving the fatigue properties may be lost. It is desirable to make the above.
熱間圧延工程終了後は必要に応じて酸洗し、その後インラインまたはオフラインで圧下率１０％以下のスキンパスまたは圧下率４０％程度までの冷間圧延を施しても構わない。 After completion of the hot rolling step, pickling may be performed if necessary, and thereafter, a skin pass with a rolling reduction of 10% or less or cold rolling to a rolling reduction of about 40% may be performed in-line or off-line.
次に、冷延鋼板として最終製品にする場合であるが、熱間での仕上げ圧延条件は特に限定しない。また、仕上げ圧延の最終パス温度（ＦＴ）はＡｒ3変態点温度未満で終了しても差し支えないが、その場合は、圧延前もしくは圧延中に強い加工組織が残留するため、続く巻取処理または加熱処理により回復、再結晶させることが望ましい。続く酸洗後の冷間圧延工程は特に限定することなく本発明の効果が得られる。 Next, there is a case where the final product is formed as a cold-rolled steel sheet, but the hot finish rolling conditions are not particularly limited. Further, the final pass temperature (FT) of the finish rolling may be completed at a temperature lower than the Ar 3 transformation point temperature, but in this case, a strong work structure remains before or during the rolling, so that the subsequent winding process or It is desirable to recover and recrystallize by heat treatment. The effect of the present invention can be obtained without any particular limitation on the cold rolling step after pickling.
この様に冷間圧延された鋼板の熱処理は連続焼鈍工程を前提としている。まず、８００℃以上の温度域で５〜１５０秒間行う。この熱処理温度が８００℃未満の場合には後の冷却においてバーリング加工性にとって好ましいベイニティックなフェライト、またはフェライトおよびベイナイトが得られない懸念があるので、熱処理温度は８００℃以上とする。また、熱処理温度の上限は特に定めないが、連続焼鈍設備の制約上実質的に９００℃以下である。 The heat treatment of the cold-rolled steel sheet is based on a continuous annealing process. First, it is performed in a temperature range of 800 ° C. or more for 5 to 150 seconds. If the heat treatment temperature is lower than 800 ° C., there is a concern that bainitic ferrite or ferrite and bainite which are preferable for burring workability may not be obtained in the subsequent cooling, so the heat treatment temperature is set to 800 ° C. or higher. Although the upper limit of the heat treatment temperature is not particularly defined, it is substantially 900 ° C. or less due to the restriction of the continuous annealing equipment.
一方、この温度域での保持時間は、５秒未満では、ＴｉおよびＮｂの炭窒化物が完全に再固溶するのに不十分であり、１５０秒超の熱処理を行ってもその効果が飽和するばかりでなく生産性を低下させるので、保持時間は５〜１５０秒間とする。 On the other hand, if the holding time in this temperature range is less than 5 seconds, the carbonitride of Ti and Nb is insufficient to completely re-dissolve solid solution, and even if heat treatment for more than 150 seconds is performed, the effect is saturated. In addition, the holding time is set to 5 to 150 seconds, since this not only reduces the productivity but also decreases the productivity.
次に冷却終了までの平均冷却速度であるが、５０℃／秒以上が必要である。これは冷却終了までの平均冷却速度が５０℃／秒未満であるとバーリング加工性にとって好ましいベイニティックなフェライト、またはフェライトおよびベイナイトの体積分率が減少してします恐れがあるからである。また、冷却速度の上限は実際の工場設備能力等を考慮すると２００℃／秒以下である。 Next, the average cooling rate up to the end of cooling is required to be 50 ° C./sec or more. This is because, if the average cooling rate until the end of cooling is less than 50 ° C./sec, bainitic ferrite or ferrite and bainite, which are preferable for burring workability, may be reduced in volume fraction. Further, the upper limit of the cooling rate is 200 ° C./sec or less in consideration of the actual factory equipment capacity and the like.
冷却終了温度は７００℃以下の温度域であることが必要であるが、連続焼鈍設備を用いる場合、冷却終了温度が５５０℃超になることは通常はないので特に配慮する必要はない。また、冷却終了温度の下限は本発明の効果を得るためには特に定める必要はない。 The cooling end temperature needs to be in a temperature range of 700 ° C. or lower. However, when using continuous annealing equipment, it is not usually necessary to take care because the cooling end temperature does not usually exceed 550 ° C. The lower limit of the cooling end temperature does not need to be particularly determined in order to obtain the effects of the present invention.
さらにその後、必要に応じてスキンパス圧延を施してもよい。 Thereafter, skin pass rolling may be performed as necessary.
酸洗後の熱延鋼板、もしくは上記の熱処理工程終了後の冷延鋼板に亜鉛めっきを施すためには、亜鉛めっき浴中に浸積し、必要に応じて合金化処理してもよい。 In order to apply galvanization to the hot-rolled steel sheet after pickling or the cold-rolled steel sheet after the heat treatment step described above, the steel sheet may be immersed in a galvanizing bath and subjected to an alloying treatment as necessary.
以下に、実施例により本発明をさらに説明する。 Hereinafter, the present invention will be further described with reference to examples.
表１に示す化学成分を有するＡ〜Ｍの鋼は、転炉にて溶製して、連続鋳造後、表２に示す加熱温度で再加熱し、粗圧延に続く仕上げ圧延で１．２〜５．５ｍｍの板厚にした後に巻き取った。ただし、表中の化学組成についての表示は質量％である。なお、表２に示すように一部については熱間圧延工程後、酸洗、冷延、熱処理を行った。板厚は０．７〜２．３ｍｍである。一方、上記鋼板のうち鋼Ｈおよび鋼Ｃ−７については、亜鉛めっきを施した。 The steels of A to M having the chemical components shown in Table 1 were melted in a converter, continuously cast, reheated at the heating temperature shown in Table 2, and subjected to rough rolling and finish rolling followed by 1.2 to 1.2 mm. After the thickness was 5.5 mm, the film was wound. However, the indication of the chemical composition in the table is% by mass. In addition, as shown in Table 2, pickling, cold rolling and heat treatment were performed on a part after the hot rolling step. The plate thickness is 0.7 to 2.3 mm. On the other hand, among the above steel sheets, steel H and steel C-7 were galvanized.
製造条件の詳細を表２に示す。ここで、「ＳＲＴ」はスラブ加熱温度、「ＦＴ」は最終パス仕上げ圧延温度、「開始時間」とは圧延終了から冷却開始までの時間、「冷却速度」とは、冷却開始から冷却停止までの平均冷却速度、「ＣＴ」は巻き取り温度である。ただし、後に冷延工程にて圧延を行う場合はこのような制限の限りではないので「―」とした。 Table 2 shows details of the manufacturing conditions. Here, “SRT” is the slab heating temperature, “FT” is the final pass finishing rolling temperature, “start time” is the time from the end of rolling to the start of cooling, and “cooling rate” is the time from the start of cooling to the stop of cooling. The average cooling rate, "CT", is the winding temperature. However, when rolling is performed later in the cold rolling process, such a limitation is not applied, so "-" is used.
このようにして得られた熱延板の引張試験は、供試材を、まず、ＪＩＳ Ｚ ２２０１記載の５号試験片に加工し、ＪＩＳ Ｚ ２２４１記載の試験方法に従って行った。試験片の形状及び寸法は、図３（ａ）、（ｂ）に示すとおりであり、鋼板１、２の継目４を溶接して溶接金属３を形成し、端部に補助板５、６を装着して試験片とした。表２に降伏強度（ＹＰ）、引張強度（ＴＳ）、破断伸び（Ｅｌ）を示す。一方、バーリング加工性（穴拡げ性）については日本鉄鋼連盟規格ＪＦＳ Ｔ １００１−１９９６記載の穴拡げ試験方法に従って評価した。表２に穴拡げ率（λ）を示す。ここで、フェライト、ベイナイト、残留オーステナイト、パーライト、マルテンサイトの体積分率とは鋼板板幅の１／４Ｗもしくは３／４Ｗ位置より切出した試料を圧延方向断面に研磨、エッチングし、光学顕微鏡を用い２００〜５００倍の倍率で観察された板厚の１／４ｔにおけるミクロ組織の面積分率で定義される。さらに、図３に示す溶接継手引張試験片にてＪＩＳ Ｚ ２２４１に準じた方法で引っ張り試験を実施し、その破断個所を目視外観観察より母材部／溶接部と分類した。継手強度の観点からこの溶接破断部は溶接部より母材部の方がより望ましい。
In the tensile test of the hot-rolled sheet obtained in this manner, the test material was first processed into a No. 5 test piece described in JIS Z 2201, and was subjected to a test method described in JIS Z 2241. The shapes and dimensions of the test pieces are as shown in FIGS. 3 (a) and 3 (b). The
なお、アーク溶接の溶接熱影響部の硬度測定はについては、ＪＩＳ Ｚ ３１０１記載の１号試験片にて、ＪＩＳ Ｚ ２２４４記載の試験方法に準じて測定した。ただし、アーク溶接は、シールドガス：ＣＯ2、ワイヤ：日鐵溶接工業（株）製ＹＭ−２８φ１．２ｍｍ、ＹＭ−６０Ｃφ１．２ｍｍ、ＹＭ−８０Ｃφ１．２ｍｍを必要に応じて使い分け、溶接速度：１００ｃｍ／分、溶接電流：２６０±１０Ａ、溶接電圧：２６±１Ｖ、供試材の板厚は研磨を行い２．６ｍｍとし、硬度測定位置は、表面より０．２５ｍｍ、測定間隔は、０．５ｍｍで、試験力は９８Ｎとした。 In addition, about the hardness measurement of the welding heat affected zone of arc welding, it measured according to the test method of JISZ2244 using the 1st test piece of JISZ3101. However, for arc welding, shield gas: CO 2 , wire: Nitto Welding Industry Co., Ltd. YM-28φ1.2mm, YM-60Cφ1.2mm, YM-80Cφ1.2mm are used as necessary, and welding speed: 100cm. / Min, welding current: 260 ± 10 A, welding voltage: 26 ± 1 V, the thickness of the test material was polished to 2.6 mm, the hardness measurement position was 0.25 mm from the surface, and the measurement interval was 0.5 mm. And the test force was 98 N.
本発明に沿うものは、鋼Ａ、Ｂ、Ｃ−１、Ｃ−７、Ｆ、Ｈ、Ｋ、Ｌ、Ｍの９鋼であり、所定の量の鋼成分を含有し、そのミクロ組織が、フェライト、またはフェライトおよびベイナイトからなることを特徴とする溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板が得られており、従って、本発明記載の方法によって評価した従来鋼の熱影響部軟化度ΔＨｖが５０以上であるのに対して有意差が認められる。さらに、鋼ＦについてはＢ添加の効果により、溶接熱影響部のうちα→γ→α変態が起こる熱履歴を受ける部位において焼入れ性が向上した結果、破断位置が母材部となっている。 According to the present invention, there are 9 steels of steels A, B, C-1, C-7, F, H, K, L, and M, each containing a predetermined amount of a steel component, and having a microstructure thereof. A burring high-strength steel sheet excellent in softening resistance of a weld heat-affected zone characterized by being composed of ferrite or ferrite and bainite has been obtained. A significant difference is recognized when the partial softening degree ΔHv is 50 or more. Further, with respect to steel F, due to the effect of the addition of B, the quenching property is improved in a part of the heat affected zone of the weld that undergoes a heat history in which α → γ → α transformation occurs. As a result, the fracture position is the base metal part.
上記以外の鋼は、以下の理由によって本発明の範囲外である。すなわち、鋼Ｃ−２は、仕上圧延終了温度（ＦＴ）が本発明請求項８の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−３は、仕上圧延終了から冷却開始までの時間が本発明請求項８の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−４は、平均冷却速度が本発明請求項８の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−５は、冷却終了温度および巻き取り温度が本発明請求項８の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−６は、巻き取り温度が本発明請求項８の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−８は、熱処理温度が本発明請求項９の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｃ−９は、保持時間が本発明請求項９の範囲外であるので、請求項１記載の目的とするミクロ組織が得られず十分な穴拡げ性（λ）が得られていない。鋼Ｄは、Ｃ*が本発明請求項１または２の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。鋼Ｅは、Ｃ*が本発明請求項１または２の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。鋼Ｅは、Ｃ添加量およびＣ*が本発明請求項１または２の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。鋼Ｇは、Ｍｏ＋Ｃｒ量が本発明請求項１の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。鋼Ｉは、Ｍｏ＋Ｃｒ量が本発明請求項１の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。鋼Ｊは、Ｃ*が本発明請求項１または２の範囲外であるので、熱影響部の軟化度（ΔＨｖ）が大きい。なお、ここに記載のＣ*は、前記した如く、Ｃ*＝Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）を意味する。
Steels other than the above are outside the scope of the present invention for the following reasons. That is, since the finish rolling temperature (FT) of the steel C-2 is out of the range of the eighth aspect of the present invention, the desired microstructure described in the first aspect cannot be obtained and sufficient hole expandability (λ) Is not obtained. Since the time from the end of finish rolling to the start of cooling of steel C-3 is outside the scope of claim 8 of the present invention, the desired microstructure described in claim 1 cannot be obtained and sufficient hole expandability (λ) Is not obtained. Since steel C-4 has an average cooling rate outside the scope of claim 8 of the present invention, the desired microstructure described in claim 1 cannot be obtained, and sufficient hole expandability (λ) cannot be obtained. Since steel C-5 has a cooling end temperature and a winding temperature outside the scope of claim 8 of the present invention, the desired microstructure described in claim 1 cannot be obtained, and sufficient hole expandability (λ) can be obtained. Not been. Since the winding temperature of steel C-6 is outside the scope of claim 8 of the present invention, the desired microstructure described in claim 1 cannot be obtained and sufficient hole expandability (λ) cannot be obtained. Since the heat treatment temperature of steel C-8 is outside the scope of claim 9 of the present invention, the desired microstructure described in claim 1 cannot be obtained, and sufficient hole expandability (λ) cannot be obtained. Since the holding time of steel C-9 is outside the scope of claim 9 of the present invention, the desired microstructure described in claim 1 cannot be obtained and sufficient hole expandability (λ) cannot be obtained. Steel C has a high degree of softening (ΔHv) of the heat-affected zone because C * is outside the scope of
１、２ 鋼板
３ 溶接金属
４ 継目
５、６ 補助板
1, 2
Claims (12)
Ｃ ：０．０１〜０．１％、
Ｓｉ：０．０１〜２％、
Ｍｎ：０．０５〜３％、
Ｐ ≦０．１％、
Ｓ ≦０．０３％、
Ａｌ：０．００５〜１％、
Ｎ ：０．０００５〜０．００５％、
Ｔｉ：０．０５〜０．５％、
を含み、さらに
０％＜Ｃ−（１２／４８Ｔｉ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％、
さらに
Ｍｏ＋Ｃｒ≧０．２％、かつＣｒ≦０．５％、Ｍｏ≦０．５％、
を満たす範囲でＣ、Ｓ、Ｎ、Ｔｉ、Ｃｒ、Ｍｏを含有し残部がＦｅ及び不可避的不純物からなる鋼であって、そのミクロ組織が、フェライト、またはフェライトおよびベイナイトからなることを特徴とする溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 In mass%,
C: 0.01-0.1%,
Si: 0.01 to 2%,
Mn: 0.05-3%,
P ≦ 0.1%,
S ≦ 0.03%,
Al: 0.005 to 1%,
N: 0.0005 to 0.005%,
Ti: 0.05-0.5%,
And 0% <C- (12 / 48Ti-12 / 14N-12 / 32S) ≦ 0.05%,
Further, Mo + Cr ≧ 0.2%, Cr ≦ 0.5%, Mo ≦ 0.5%,
A steel containing C, S, N, Ti, Cr, Mo in the range satisfying the condition, and the balance being Fe and unavoidable impurities, wherein the microstructure is made of ferrite or ferrite and bainite. Burring high strength steel sheet with excellent softening resistance in the weld heat affected zone.
Ｎｂ：０．０１〜０．５％、
を含み、さらに
０＜Ｃ−（１２／４８Ｔｉ＋１２／９３Ｎｂ−１２／１４Ｎ−１２／３２Ｓ）≦０．０５％、
を満たす範囲でＮｂを含有し残部がＦｅ及び不可避的不純物からなる鋼であることを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 The steel further comprises, in mass%:
Nb: 0.01-0.5%,
And 0 <C- (12 / 48Ti + 12 / 93Nb-12 / 14N-12 / 32S) ≦ 0.05%,
A high-strength burring steel sheet excellent in softening resistance of a weld heat-affected zone, wherein Nb is a steel containing Nb in a range satisfying the following, and the balance is Fe and unavoidable impurities.
Ｃａ：０．０００５〜０．００２％、
ＲＥＭ：０．０００５〜０．０２％
の一種または二種を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 The steel according to claim 1 or 2, further comprising:
Ca: 0.0005 to 0.002%,
REM: 0.0005-0.02%
A high-strength burring steel sheet having excellent resistance to softening of the heat affected zone by welding, characterized by containing one or two of the following.
Ｃｕ：０．２〜１．２％
を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 The steel according to any one of claims 1 to 3, further comprising:
Cu: 0.2-1.2%
A high-strength burring steel sheet having excellent resistance to softening of the heat affected zone by welding.
Ｎｉ：０．１〜０．６％
を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 The steel according to any one of claims 1 to 4, further comprising:
Ni: 0.1 to 0.6%
A high-strength burring steel sheet having excellent resistance to softening of the heat affected zone by welding.
Ｂ ：０．０００２〜０．００２％
を含有することを特徴とする、溶接熱影響部の耐軟化性に優れたバーリング性高強度鋼板。 The steel according to any one of claims 1 to 5, further comprising:
B: 0.0002 to 0.002%
A high-strength burring steel sheet having excellent resistance to softening of the heat affected zone by welding.
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| WO2013022043A1 (en) * | 2011-08-09 | 2013-02-14 | 新日鐵住金株式会社 | Hot-rolled steel sheet having high yield ratio and excellent low-temperature impact energy absorption and haz softening resistance and method for producing same |
| WO2014171427A1 (en) | 2013-04-15 | 2014-10-23 | 新日鐵住金株式会社 | Hot-rolled steel sheet |
| WO2016005780A1 (en) * | 2014-07-11 | 2016-01-14 | Arcelormittal Investigación Y Desarrollo Sl | Hot-rolled steel sheet and associated manufacturing method |
| KR20170093886A (en) | 2015-02-20 | 2017-08-16 | 신닛테츠스미킨 카부시키카이샤 | Hot-rolled steel sheet |
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