JP4291711B2 - High burring hot rolled steel sheet having bake hardenability and method for producing the same - Google Patents
High burring hot rolled steel sheet having bake hardenability and method for producing the same Download PDFInfo
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Description
本発明は、焼付け硬化性を有する高バーリング熱延鋼板およびその製造方法に関するものであり、特に４９０ＭＰａ級超の引張強度の鋼板であっても優れたバーリング加工性を発現させる均一なミクロ組織を有し、厳しい伸びフランジ加工が要求される部品でも容易に成形できるだけでなく、プレスによるひずみ導入と塗装焼き付け処理により１グレード上の鋼板を適用した場合の設計強度に相当するプレス品強度を得ることができる。 The present invention relates to a high burring hot-rolled steel sheet having bake hardenability and a method for producing the same, and in particular, has a uniform microstructure that exhibits excellent burring workability even for a steel sheet having a tensile strength exceeding 490 MPa. In addition, not only parts that require strict stretch flange processing can be easily formed, but also can obtain a pressed product strength equivalent to the design strength when one grade steel plate is applied by introducing strain and painting and baking by press. it can.
自動車の燃費向上などのために車体軽量化を目的としたゲージダウンへの要求は近年益々高まっており、ゲージダウンを前提に如何にしてプレス品強度特性を維持するかが車体軽量化の課題である。そこで、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が、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いる。従ってより安価かつ広い範囲に自動車の軽量化を推進するためには、鋼板適用によるゲージダウンの実現が必要とされている。 The demand for gauge down for the purpose of reducing the weight of the car body for the purpose of improving the fuel efficiency of automobiles has been increasing in recent years. is there. Therefore, although light metals such as Al alloys and high strength steel plates are being applied to automobile members, although light metals such as Al alloys have the advantage of high specific strength, they are significantly more expensive than steel. Its application is limited to special uses. Therefore, in order to promote the weight reduction of automobiles at a lower cost and in a wider range, it is necessary to realize gauge down by applying a steel plate.
鋼板を用いたゲージダウンの手段として高強度鋼板の適用が検討されている。しかし、材料の高強度化は一般的に成形性（加工性）等の材料特性を劣化させるため、材料特性を劣化させずに如何に高強度化を図るかが高強度鋼板開発のカギになる。特に内板部材、構造部材、足廻り部材用鋼板に求められる特性としては、伸びフランジ性（バーリング性）、延性、疲労耐久性および耐食性等が重要であり、高強度とこれら特性を如何に高次元でバランスさせるかが重要である。 Application of high-strength steel sheets is being studied as a means of gauge down using steel sheets. However, increasing the strength of materials generally degrades material properties such as formability (workability), so how to increase strength without deteriorating material properties is the key to developing high-strength steel sheets. . In particular, stretch flangeability (burring properties), ductility, fatigue durability, and corrosion resistance are important properties required for steel sheets for inner plate members, structural members, and suspension members, and how high these strengths are. It is important to balance in dimension.
このように高強度化と諸特性、特に成形性を両立するために、例えば鋼のミクロ組織中に残留オーステナイトを含むことで、成形中にＴＲＩＰ（TRansformation Induced Plasticity ）現象を発現させることで飛躍的に成形性（延性および深絞り性）を向上させたＴＲＩＰ鋼が開示されている（例えば特許文献１、２）。しかし、上記技術はバーリング性を向上させることを念頭に置いたものではないため、サスペンションアーム等をはじめとする足廻り部材のように厳しいバーリング加工が施される部品に適用するには、バーリング加工性が不十分である。
In this way, in order to achieve both high strength and various properties, especially formability, for example, by including residual austenite in the microstructure of steel, the TRIP (TRansformation Induced Plasticity) phenomenon is manifested during molding. Discloses TRIP steel having improved formability (ductility and deep drawability) (for example,
一方、高強度とバーリング加工性を両立する手段として、例えば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を得る発明が開示されている（例えば特許文献３）。さらに、Ｔｉ，Ｎｂを添加することにより第二相を低減して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を得る発明が開示されている（例えば特許文献４）。 On the other hand, as a means for achieving both high strength and burring workability, for example, Ti and Nb are added to reduce the second phase, and TiC and NbC are precipitated and strengthened in polygonal ferrite, which is the main phase. An invention for obtaining a high-strength hot-rolled steel sheet having excellent flangeability is disclosed (for example, Patent Document 3). Furthermore, an invention for obtaining a high-strength hot-rolled steel sheet excellent in stretch flangeability by reducing the second phase by adding Ti and Nb to make the microstructure as acicular ferrite and strengthening by precipitation with TiC and NbC is disclosed. (For example, Patent Document 4).
しかしながら、これら鋼板はバーリング加工性には優れるものの、同一強度の上記ＴＲＩＰ鋼等と比較して延性および深絞り性が著しく劣っており、絞り成形が用いられるセンターピラーリインフォース等の構造部材への適用は難しい。すなわち、強度、延性、バーリング加工性はそれぞれトレードオフ関係にあり、これら全てを高次元でバランスさせることは非常に難しいのが現状である。 However, although these steel plates are excellent in burring workability, they are significantly inferior in ductility and deep drawability compared to the above-mentioned TRIP steel etc. of the same strength, and are applied to structural members such as center pillar reinforcements in which draw forming is used. Is difficult. That is, strength, ductility, and burring workability are in a trade-off relationship, and it is very difficult to balance all of these at a high level.
このような課題を解決する手段として、プレス成形時には強度が低く、プレスによるひずみの導入と後の塗装焼き付け処理にてプレス品の強度を向上させるＢＨ（Bake Hardening）鋼板が提案されている。すなわち、比較的強度の低い状態でプレス加工を行うことで成形性を担保し、焼付け硬化によりプレス品強度を確保するということである。
ＢＨ性を向上させるためには固溶ＣやＮの増加させることが有効であるが、一方でこれら固溶元素の増加は常温での時効劣化を悪化させるため、ＢＨ性と耐常温時効劣化を両立させることが重要な技術となる。
As means for solving such a problem, a BH (Bake Hardening) steel sheet has been proposed that has low strength during press forming and improves the strength of the pressed product by introducing strain due to pressing and subsequent coating baking treatment. That is, by performing the pressing process in a relatively low strength state, the moldability is ensured and the strength of the pressed product is ensured by baking and curing.
In order to improve the BH property, it is effective to increase the solute C and N. On the other hand, the increase of these solute elements deteriorates the aging deterioration at room temperature, so that the BH property and the room temperature aging deterioration are reduced. It is an important technology to achieve both.
以上のような必要性から、固溶Ｎの増加によりＢＨ性を向上させ、結晶粒細粒化により増加した粒界面積の効果で常温における固溶Ｃ，Ｎの拡散を抑制することでＢＨ性と耐常温時効劣化を両立させる技術が開示されている（例えば特許文献５）。しかしながら、ミクロ組織がフェライト−パーライト組織であり、４９０ＭＰａ級超の強度を得ことは難しく、さらなるゲージダウンには対応できない。 From the above necessity, the BH property is improved by improving the BH property by increasing the solid solution N and suppressing the diffusion of the solid solution C and N at room temperature by the effect of the grain boundary area increased by the grain refinement. And a technique that achieves both room temperature and aging resistance (for example, Patent Document 5). However, the microstructure is a ferrite-pearlite structure, and it is difficult to obtain strength exceeding 490 MPa class, and it cannot cope with further gauge down.
一方、ミクロ組織中にマルテンサイト等の硬質第二相を含有させることにより、高強度でありながらＢＨ性を有する鋼板を製造する技術が開示されている（例えば特許文献６）。しかしながら、このようにマルテンサイト等の硬質第二相を含むミクロ組織をもつ鋼板は、元来バーリング加工性が著しく低いことが知られており、対象とする自動車部品への適用は困難である。 On the other hand, a technique for manufacturing a steel sheet having high strength and BH properties by including a hard second phase such as martensite in the microstructure is disclosed (for example, Patent Document 6). However, it is known that a steel sheet having a microstructure including a hard second phase such as martensite is inherently extremely low in burring workability, and it is difficult to apply to a target automobile part.
さらに、前記特許文献１、２のＴＲＩＰ鋼は優れたＢＨ性を示すものも存在するが、上記特許文献６の鋼板と同じく複合組織鋼であるがゆえにバーリング加工性が低い。また、前記特許文献３、４の鋼板はＴｉ，Ｎｂの析出強化を適用しているため、ＢＨ性に必要な固溶Ｃ，ＮがｓｃａｖｅｎｇｉｎｇされてしまいＢＨ性が期待できない。
そこで本発明は、優れたバーリング加工性を有するとともに４９０ＭＰａ級超の強度で５０ＭＰａ以上のＢＨ量を得られる、焼付け硬化性を有する高バーリング熱延鋼板およびその製造方法を提供する。すなわち本発明は、優れたバーリング加工性を発現させる均一なミクロ組織を有し、プレスによるひずみ導入と塗装焼き付け処理により、１グレード上の鋼板を適用した場合の設計強度に相当するプレス品強度を得ることができる、焼付け硬化性を有する高バーリング熱延鋼板、およびその鋼板を安価に安定して製造できる方法を提供することを目的とする。 Therefore, the present invention provides a high burring hot-rolled steel sheet having bake hardenability, which has excellent burring workability and can obtain a BH amount of 50 MPa or more with strength exceeding 490 MPa class and a method for producing the same. That is, the present invention has a uniform microstructure that expresses excellent burring workability, and has a press product strength equivalent to the design strength when a steel plate of one grade is applied by strain introduction by press and paint baking treatment. It is an object of the present invention to provide a high burring hot-rolled steel sheet having bake hardenability and a method capable of stably and inexpensively manufacturing the steel sheet.
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４９０ＭＰａ級超の鋼板の製造プロセスを念頭において、焼付け硬化性を有する高バーリング熱延鋼板を得るべく鋭意研究を重ねた結果、以下の手段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 inventors of the present invention are eager to obtain a high burring hot-rolled steel sheet having bake hardenability in consideration of the manufacturing process of a steel sheet exceeding 490 MPa class currently produced on an industrial scale by a production facility that is currently normally employed. As a result of repeated research, the present inventors have newly found that the following means are very effective.
即ち、本発明の要旨は以下の通りである。
（１） 質量％にて、
Ｃ ：０．０１〜０．１％、 Ｓｉ：０．００５〜２％、
Ｍｎ：０．１〜３％、 Ｐ ≦０．１％、
Ｓ ≦０．０３％、 Ａｌ：０．００１〜０．１％、
Ｎ ≦０．００５％、 Ｔｉ：０．０５〜０．２％
を含み、さらに
Ｔｉ（＊）＝Ｔｉ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％
かつ
Ｃ−（１２／４８）Ｔｉ（＊）≦０．０５％
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板であって、そのミクロ組織が主にベイニティックフェライトであり、旧オーステナイト粒の平均粒径が８μｍ超〜８０μｍであることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。
（２） 前記（１）に記載の鋼が、さらに質量％にて、
Ｎｂ：０．０１〜０．２％
を含み、さらに
Ｔｉ（＊）＝Ｔｉ＋（４８／９３）Ｎｂ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％かつ
Ｃ−（１２／４８）Ｔｉ（＊）≦０．０５％
を満たす範囲でＮｂを含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。
（３） 前記（１）または（２）に記載の鋼が、さらに質量％にて、
Ｃａ：０．０００５〜０．００５％、
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焼付け硬化性を有する高バーリング熱延鋼板。
（４） 前記（１）〜（３）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焼付け硬化性を有する高バーリング熱延鋼板。
（５） 前記（１）〜（４）のいずれか１項に記載の鋼が、さらに質量％にて、
Ｍｏ：０．０５〜１％、 Ｖ ：０．０２〜０．２％、
Ｃｒ：０．０１〜１％
の一種または二種以上を含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。
（６） 前記（１）〜（５）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％、 Ｎｉ：０．１〜０．６％
の一種または二種を含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。
（７） 前記（１）〜（３）のいずれか１項に記載の薄鋼板に亜鉛めっきが施されていることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。
That is, the gist of the present invention is as follows.
(1) In mass%
C: 0.01 to 0.1%, Si: 0.005 to 2%,
Mn: 0.1 to 3%, P ≦ 0.1%,
S ≦ 0.03%, Al: 0.001 to 0.1%,
N ≦ 0.005%, Ti: 0.05 to 0.2%
Ti (*) = Ti− (48/14) N− (48/32) S ≧ 0%
And C- (12/48) Ti (*) ≦ 0.05%
Is a steel plate that contains C, S, N, Ti in the range satisfying the following, the balance is Fe and inevitable impurities, the microstructure is mainly bainitic ferrite, the average grain size of the prior austenite grains is A high burring hot-rolled steel sheet having bake hardenability, characterized by being over 8 μm to 80 μm.
(2) The steel according to (1) is further in mass%,
Nb: 0.01 to 0.2%
Ti (*) = Ti + (48/93) Nb− (48/14) N− (48/32) S ≧ 0% and C− (12/48) Ti (*) ≦ 0.05%
A high burring hot-rolled steel sheet having bake hardenability characterized by containing Nb in a range satisfying the above.
(3) The steel according to (1) or (2) is further in mass%,
Ca: 0.0005 to 0.005%,
REM: 0.0005 to 0.02%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or two of the following.
(4) The steel according to any one of (1) to (3) is further in mass%,
B: 0.0002 to 0.002%
A high burring hot-rolled steel sheet having bake hardenability, comprising:
(5) The steel according to any one of (1) to (4), further in mass%,
Mo: 0.05 to 1%, V: 0.02 to 0.2%,
Cr: 0.01 to 1%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or more of the above.
(6) The steel according to any one of (1) to (5), further in mass%,
Cu: 0.2-1.2%, Ni: 0.1-0.6%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or two of the following.
(7) A high burring hot-rolled steel sheet having bake hardenability, wherein the thin steel sheet according to any one of (1) to (3) is galvanized.
（８） 前記（１）〜（６）のいずれか１項に記載の成分を有する薄鋼板を得るために熱間圧延をする際に、該成分を有する鋼片を粗圧延後に最終段の圧下率が１〜１５％の仕上圧延をＡｒ3 変態点温度＋５０℃以上の温度域で終了後、０．５〜５秒間で冷却を開始し、Ａｒ3 〜Ｂｆの温度域を８０〜５００℃／sec の冷却速度で冷却し、さらに５００℃以下まで２０℃／sec 以上の冷却速度で冷却した後、５００℃以下で巻き取ることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板の製造方法。
（９） 前記（８）に記載の熱間圧延に際し、仕上圧延開始温度を１０００℃以上とすることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板の製造方法。
（１０） 前記（８）または（９）に記載の熱間圧延に際し、鋼片を粗圧延終了した後の粗バーを、仕上圧延開始までの間および粗バーの仕上圧延中の、いずれか一方または両方で加熱することを特徴とする、焼付け硬化性を有する高バーリング熱延鋼板の製造方法。（１１） 前記（８）〜（１０）のいずれか１項に記載の熱間圧延に際し、粗圧延終了から仕上圧延開始までの間にデスケーリングを行うことを特徴とする、焼付け硬化性を有する高バーリング熱延鋼板の製造方法。
（１２） 前記（８）〜（１１）のいずれか１項に記載の熱間圧延後、得られた熱延鋼板を亜鉛めっき浴中に浸漬させて鋼板表面を亜鉛めっきすることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板の製造方法。
（１３） 前記（１２）に記載の製造方法に際し、亜鉛めっき後、合金化処理することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板の製造方法。
(8) When hot-rolling to obtain a thin steel sheet having the component according to any one of (1) to (6), the steel slab having the component is subjected to final rolling reduction after rough rolling. After finishing rolling at a rate of 1 to 15% in the temperature range of Ar3 transformation temperature + 50 ° C or higher, cooling is started in 0.5 to 5 seconds, and the temperature range of Ar3 to Bf is 80 to 500 ° C / sec. A method for producing a high burring hot-rolled steel sheet having bake hardenability, wherein the steel sheet is cooled at a cooling rate, further cooled to 500 ° C. or less at a cooling rate of 20 ° C./sec or more, and then wound at 500 ° C. or less.
(9) A method for producing a high burring hot-rolled steel sheet having bake hardenability, wherein the finish rolling start temperature is set to 1000 ° C. or higher during the hot rolling described in (8).
(10) In the hot rolling described in (8) or (9) above, either the rough bar after the rough rolling of the steel slab is finished until the start of finish rolling or during the finish rolling of the coarse bar. Alternatively, a method for producing a high burring hot-rolled steel sheet having bake hardenability, characterized by heating both. (11) In the hot rolling described in any one of (8) to (10), the descaling is performed from the end of rough rolling to the start of finish rolling, and has bake curability. Manufacturing method of high burring hot-rolled steel sheet.
(12) After hot rolling according to any one of (8) to (11), the obtained hot rolled steel sheet is immersed in a galvanizing bath to galvanize the steel sheet surface. A method for producing a high burring hot rolled steel sheet having bake hardenability.
(13) A method for producing a high burring hot-rolled steel sheet having bake hardenability, characterized in that, in the production method according to (12), after galvanization, alloying is performed.
本発明は、焼付け硬化性を有する高バーリング熱延鋼板およびその製造方法に関するものであり、これらの鋼板を用いることにより厳しいバーリング加工が要求される部品でも容易に成形できるだけでなく、４９０ＭＰａ級超の強度で安定して５０ＭＰａ以上のＢＨ量を得られるので、プレスによるひずみ導入と塗装焼き付け処理により１グレード上の鋼板を適用した場合の設計強度に相当するプレス品強度を得ることができるため、工業的価値が高い発明であると言える。 The present invention relates to a high burring hot-rolled steel sheet having bake hardenability and a method for producing the same, and by using these steel sheets, not only parts that require severe burring processing can be easily formed, but also exceeding 490 MPa class. Since a BH amount of 50 MPa or more can be obtained stably in strength, it is possible to obtain the strength of a pressed product corresponding to the design strength when a steel sheet of one grade or higher is applied by strain introduction by press and paint baking treatment. It can be said that the invention is highly valuable.
以下に、本発明に至った基礎的研究結果について説明する。
発明者らは、バーリング加工性を維持しつつ高い焼付け硬化性（ＢＨ性）を得るために鋭意研究を重ねた結果、以下の結論に至った。すなわち、(1) バーリング加工性を確保するためにＴｉ等で粗大な炭化物の析出を抑制した系において、さらにバーリング加工性を向上させるためには粒界脆化元素であるＰ等の局在化を押さえることが重要である。さらに、(2) このような固溶ＣをＴｉ等の析出物として固定してしまうような化学量論組成に近い成分であっても、製造条件により十分なＢＨ量を確保できる。
Hereinafter, the basic research results that led to the present invention will be described.
As a result of intensive studies to obtain high bake hardenability (BH property) while maintaining burring workability, the inventors have reached the following conclusion. (1) In a system in which precipitation of coarse carbides with Ti or the like is suppressed to ensure burring workability, in order to further improve burring workability, localization of P or the like as an intergranular embrittlement element It is important to hold down. Furthermore, (2) even if it is a component close to the stoichiometric composition that fixes such solute C as a precipitate such as Ti, a sufficient amount of BH can be secured depending on the production conditions.
そこで発明者らは上記のような新しい知見に基づいて、ＢＨ性とバーリング加工性を両立できるような製造条件およびミクロ組織との関係を、Ｔｉ炭化物の析出および粒界脆化元素であるＰ等の局在化という現象に密接に関係していると思われる旧オーステナイト粒径という観点で調査するために、次のような実験を行った。表１に示す鋼成分の鋳片を溶製し様々な製造プロセスで製造した２ｍｍ厚の鋼板を準備し、それらについてＢＨ性とバーリング加工性およびミクロ組織を調査した。 Therefore, based on the above-described new findings, the inventors have established the relationship between the manufacturing conditions and the microstructure that can achieve both BH properties and burring workability, precipitation of Ti carbides, P that is an intergranular embrittlement element, and the like. In order to investigate from the viewpoint of the prior austenite grain size, which seems to be closely related to the phenomenon of localization of the following, the following experiment was conducted. 2 mm-thick steel plates prepared by melting various slabs of steel components shown in Table 1 and prepared by various manufacturing processes were prepared, and BH properties, burring workability, and microstructures were investigated.
ＢＨ性は以下の手順に従い評価した。それぞれの鋼板よりＪＩＳ Ｚ ２２０１に記載の５号試験片を切出し、これら試験片に２％の引張予ひずみを試験片に付与した後、１７０℃×２０分の塗装焼き付け工程相当の熱処理を施してから再度引張試験を実施した。引張試験はＪＩＳ Ｚ ２２４１の方法に従った。ここでＢＨ量とは、再引張での上降伏点から２％の引張り予ひずみの流動応力を差し引いたと定義される。
バーリング加工性は、日本鉄鋼連盟規格ＪＦＳ Ｔ １００１−１９９６記載の穴拡げ試験方法に従い、穴拡げ値にて評価した。
BH property was evaluated according to the following procedure. Cut out No. 5 test pieces described in JIS Z 2201 from each steel plate, and after applying 2% tensile pre-strain to the test pieces, heat treatment equivalent to a coating baking process of 170 ° C. × 20 minutes was performed. The tensile test was performed again. The tensile test followed the method of JIS Z 2241. Here, the BH amount is defined as subtracting 2% tensile prestrained flow stress from the upper yield point in re-tensioning.
Burring workability was evaluated by a hole expansion value according to a hole expansion test method described in Japan Iron and Steel Federation Standard JFS T 1001-1996.
一方、ミクロ組織の調査は鋼板板幅の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表層下０．２ｍｍ、板厚の１／４ｔ、１／２ｔにおける視野の写真にて行った。ミクロ組織の体積分率とは、上記金属組織写真において面積分率で定義される。
ここでベイニティックフェライトとは、日本鉄鋼協会基礎研究会ベイナイト調査研究部会編、「低炭素鋼のベイナイト組織と変態挙動に関する最近の研究」、ベイナイト調査研究部会最終報告書（１９９４年、日本鉄鋼協会）に記載されているように、無拡散でせん断的機構により生成する変態組織と定義されるミクロ組織である。
On the other hand, the microstructure is examined by grinding a sample cut from a 1/4 W or 3/4 W position of the steel plate width to a cross section in the rolling direction, etching using a Nital reagent, and 200-500 times magnification using an optical microscope. The observation was carried out with photographs of the field of view at 0.2 mm below the surface layer and at 1/4 t and 1/2 t of the plate thickness. The volume fraction of the microstructure is defined as an area fraction in the metal structure photograph.
Here, bainitic ferrite refers to the latest research on the bainite structure and transformation behavior of low-carbon steel, edited by the Japan Iron and Steel Institute Basic Research Group, Bainite Research Group, 1994, Japan Steel This is a microstructure defined as a metamorphic structure generated by a non-diffusion and shearing mechanism, as described in the Association).
次に旧オーステナイト粒の平均粒径の測定であるが、ナイタール試薬にてエッチングし光学顕微鏡にて観察した同一試料を再度研磨し、特開平０６−２０７２７９号公報に記載の腐食液および腐食方法に従いエッチングした後、ＪＩＳ Ｇ ０５５２記載の切断法を用い、その測定値より求めた粒度番号Ｇより、断面積１ｍｍ2 当たりの結晶粒の数ｍをｍ＝８×２G より求め、このｍよりｄm ＝１／√ｍで得られる平均粒径ｄm を平均粒径と定義する。 Next, the average particle size of prior austenite grains is measured. The same sample that has been etched with a Nital reagent and observed with an optical microscope is polished again, and in accordance with the corrosive liquid and the corrosive method described in JP-A-06-207279. After etching, using the cutting method described in JIS G 0552, the number m of crystal grains per 1 mm 2 in cross-sectional area is obtained from m = 8 × 2 G from the particle size number G obtained from the measured value. the average particle size d m obtained by m = 1 / √m defined as the average particle size.
上記の方法にてＢＨ量、穴拡げ値を測定した結果において、旧オーステナイト粒の平均粒径とＢＨ量、穴拡げ値（λ）の関係を図１に示す。ＢＨ量、穴拡げ値（λ）と旧オーステナイト粒の平均粒径には非常に強い相関があり、平均結晶粒径が８μｍ超〜８０μｍであるとＢＨ量と穴拡げ値（λ）が高い値で両立できることを新たに知見した。 FIG. 1 shows the relationship between the average grain size of the prior austenite grains, the BH amount, and the hole expansion value (λ) as a result of measuring the BH amount and the hole expansion value by the above method. There is a very strong correlation between the BH amount and the hole expansion value (λ) and the average grain size of the prior austenite grains. When the average crystal grain size is more than 8 μm to 80 μm, the BH amount and the hole expansion value (λ) are high. It was newly discovered that it can be compatible.
このメカニズムは必ずしも明らかではないが以下のように推定される。
特開平０６−２０７２７９号公報に記載のピクリン酸、ドデシルベンゼンスルフォン酸ナトリウム等よりなる腐食液にて現出される粒界は、変態前のオーステナイト粒界（旧オーステナイト粒界）と推定される。
この粒の粒径が８μｍ以下であると、フェライトの析出サイトとなるオーステナイト粒界の単位体積あたりの面積率が大きくなり、その結果、炭化物の析出サイトとなるγ→α°B 変態境界の単位体積あたりの面積率が大きくなる。また、γ→α°B 変態境界の単位体積あたり面積率が大きいと、固溶Ｃの拡散距離が短くなるため、析出サイトが多いことと相俟って炭化物の析出が促進されることとなる。するとＢＨ性に必要な固溶Ｃが炭化物として析出してしまうためＢＨ量が低下する。
Although this mechanism is not necessarily clear, it is estimated as follows.
Grain boundaries appearing in a corrosive solution composed of picric acid, sodium dodecylbenzene sulfonate, etc. described in JP-A-06-207279 are presumed to be austenite grain boundaries before transformation (former austenite grain boundaries).
When the grain size is 8 μm or less, the area ratio per unit volume of the austenite grain boundary, which becomes the ferrite precipitation site, increases, and as a result, the unit of the γ → α ° B transformation boundary that becomes the carbide precipitation site. The area ratio per volume increases. In addition, when the area ratio per unit volume of the γ → α ° B transformation boundary is large, the diffusion distance of the solid solution C is shortened, so that precipitation of carbides is promoted in combination with a large number of precipitation sites. . Then, since the solid solution C required for BH property will precipitate as a carbide | carbonized_material, the amount of BH will fall.
一方、変態前のオーステナイト粒界（旧オーステナイト粒界）が８０μｍ超では、旧オーステナイト粒界の単位体積あたりの面積率が小さくなる。その結果、Ｐ等の粒界脆化元素がオーステナイト粒界に偏析すると、オーステナイト粒界の単位体積あたりの面積率が小さいので粗大に局在化し易くなる。さらに局在化したＰ等はＣと比較して拡散速度が遅いために、γ→α°B 変態後もこの局在化は是正されない。従って、このＰ等の粒界脆化元素の粗大な局在化個所がバーリング割れの起点となるため穴拡げ性が低下する。 On the other hand, when the austenite grain boundary before transformation (former austenite grain boundary) exceeds 80 μm, the area ratio per unit volume of the prior austenite grain boundary becomes small. As a result, when grain boundary embrittlement elements such as P are segregated at the austenite grain boundaries, the area ratio per unit volume of the austenite grain boundaries is small, so that they are easily localized coarsely. Furthermore, since localized P or the like has a slower diffusion rate than C, this localization is not corrected even after the γ → α ° B transformation. Therefore, since the coarse localized portion of the grain boundary embrittlement element such as P becomes the starting point of burring crack, the hole expandability is lowered.
さらに、図１に示すように巻き取り温度によってＴｉ等によるＣの固定が促進され、ミクロ組織がポリゴナルフェライト（Ｑｕａｓｉ−ポリゴナルフェライトを含む）となる場合あるが、このミクロ組織ではバーリング割れの起点となるセメンタイト等の粗大な炭化物がほとんどないため、穴拡げ性は優れるものの、ＢＨ性を得るための固溶ＣがＴｉ等の炭化物の析出に消費されてしまい、ＢＨ性を発現させるのに十分な固溶Ｃが存在しなくなり、ＢＨ性がほとんど期待できない。従って、ミクロ組織が主にベイニティックフェライトの場合に十分なＢＨ量が得られる。なお図中のＢＦ、ＰＦはそれぞれの典型例としてその面積率が１００％の場合を示している。
本発明においては、上記で評価した２％予ひずみでのＢＨ量が優れるのみでなく、１０％予ひずみでのＢＨ量が３０ＭＰａ以上、１０％予ひずみでの引張強度の上昇代（ΔＴＳ）が３０ＭＰａ以上得られることも付記しておく。
Furthermore, as shown in FIG. 1, the fixing of C by Ti or the like is promoted by the winding temperature, and the microstructure may become polygonal ferrite (including quasi-polygonal ferrite), but in this microstructure, burring cracking occurs. Although there is almost no coarse carbide such as cementite as a starting point, the hole expandability is excellent, but the solid solution C for obtaining BH property is consumed for precipitation of carbides such as Ti, and the BH property is expressed. Sufficient solid solution C does not exist and BH property is hardly expected. Therefore, a sufficient amount of BH can be obtained when the microstructure is mainly bainitic ferrite. In addition, BF and PF in the figure show a case where the area ratio is 100% as a typical example.
In the present invention, not only the BH amount at the 2% pre-strain evaluated above is excellent, but the BH amount at the 10% pre-strain is 30 MPa or more, and the increase in tensile strength (ΔTS) at the 10% pre-strain is It is also noted that 30 MPa or more can be obtained.
続いて、本発明の化学成分の限定理由について説明する。
Ｃは、本発明において最も重要な元素の一つである。０．１％超含有していると伸びフランジ割れの起点となる炭化物が増加し、穴拡げ値が劣化するので、０．１％以下とする。延性を考慮すると０．０６％以下が望ましい。また０．０１％未満ではＢＨ量を低下させてしまう怖れがあるので、０．０１％以上とする。
Then, the reason for limitation of the chemical component of this invention is demonstrated.
C is one of the most important elements in the present invention. If it exceeds 0.1%, the carbide that becomes the starting point of stretch flange cracking increases and the hole expansion value deteriorates, so the content is made 0.1% or less. Considering ductility, 0.06% or less is desirable. Further, if it is less than 0.01%, there is a fear that the amount of BH may be reduced, so the content is made 0.01% or more.
Ｓｉは、冷却中に伸びフランジ割れの起点となる鉄炭化物の析出を抑制する効果があるので０．００５％以上添加するが、２％を超えて添加してもその効果が飽和する。従ってその上限を２％とする。さらに０．３％超では化成処理性を劣化させる恐れがあるので、望ましくはその上限を０．３％とする。 Since Si has an effect of suppressing precipitation of iron carbide that becomes the starting point of stretch flange cracking during cooling, it is added in an amount of 0.005% or more, but even if added over 2%, the effect is saturated. Therefore, the upper limit is set to 2%. Further, if it exceeds 0.3%, there is a possibility that the chemical conversion processability is deteriorated, so the upper limit is desirably set to 0.3%.
Ｍｎは、オーステナイト域温度を低温側に拡大させ圧延終了後の冷却中に、本発明の要件であるベイニティックフェライトを得やすくする効果があるので、０．１％以上添加する。しかしながら、Ｍｎは３％超添加してもその効果が飽和するので、その上限を３％とする。また、Ｍｎ以外にＳによる熱間割れの発生を抑制する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。 Mn has an effect of making it easy to obtain bainitic ferrite, which is a requirement of the present invention, by expanding the austenite temperature to the low temperature side and cooling after the end of rolling, so 0.1% or more is added. However, since the effect is saturated even if Mn is added in excess of 3%, the upper limit is made 3%. In addition, in addition to Mn, when an element that suppresses the occurrence of hot cracking due to S is not sufficiently added, it is desirable to add an amount of Mn that satisfies Mn / S ≧ 20 by mass%.
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ので、０．１％以下とする。ただし、穴拡げ性や溶接性を考慮すると０．０２％以下が望ましい。 P is an impurity and is preferably as low as possible. If contained over 0.1%, the workability and weldability are adversely affected. However, considering hole expansibility and weldability, 0.02% or less is desirable.
Ｓは、熱間圧延時の割れを引き起こすばかりでなく、多すぎると穴拡げ性を劣化させるＡ系介在物を生成する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。ただし、ある程度の穴拡げ性を必要とする場合は０．０１％以下が、さらに高い穴拡げが要求される場合は０．００３以下が望ましい。 S not only causes cracking during hot rolling, but if it is too much, it generates A-based inclusions that degrade the hole expandability, so it should be reduced as much as possible. It is. However, 0.01% or less is desirable when a certain degree of hole expansion is required, and 0.003 or less is desirable when higher hole expansion is required.
Ａｌは、溶鋼脱酸のために０．００１％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０．１％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０．０６％以下とする。 Al needs to be added in an amount of 0.001% or more for deoxidation of molten steel, but the upper limit is set to 0.1% because of an increase in cost. Moreover, when adding too much, a nonmetallic inclusion will be increased and elongation will be degraded, Therefore It is 0.06% or less desirably.
Ｔｉは、本発明における最も重要な元素の一つである。Ｔｉは析出強化により鋼板の強度上昇に寄与する。さらにセメンタイト等の粗大な炭化物の析出を抑制し、バーリング加工性を向上させる効果がある。ただし、０．０５％未満ではこの効果が不十分であり、０．２％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、０．２％以下とする。
さらにＴｉは、製造条件を限定することによりＢＨ性を発現させる固溶Ｃ量を最適化する効果がある。すなわち、Ｔｉ量が少なすぎると高温で窒化物、硫化物を形成し、ＢＨ性に必要な固溶Ｃ量が得られないので、Ｔｉ（＊）＝Ｔｉ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％とする。また、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過剰な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、さらにＢＨ量確保に必要な最適な固溶Ｃ量をえるためには、Ｃ−（１２／４８）Ｔｉ（＊）≦０．０５％、好ましくは≦０．０３％の条件を満たすことが必要である。
Ti is one of the most important elements in the present invention. Ti contributes to an increase in strength of the steel sheet by precipitation strengthening. Furthermore, it has the effect of suppressing precipitation of coarse carbides such as cementite and improving burring workability. However, if it is less than 0.05%, this effect is insufficient, and if it exceeds 0.2%, the effect is not only saturated but also the alloy cost is increased. Therefore, the Ti content is 0.05% or more and 0.2% or less.
Furthermore, Ti has the effect of optimizing the amount of solid solution C that expresses BH properties by limiting the production conditions. That is, if the amount of Ti is too small, nitrides and sulfides are formed at high temperatures, and the amount of dissolved C necessary for the BH property cannot be obtained. Therefore, Ti (*) = Ti− (48/14) N− (48 / 32) S ≧ 0%. In addition, in order to contribute to the improvement of burring workability by depositing and fixing excess C that causes carbides such as cementite which deteriorates burring workability, in order to obtain the optimum amount of solute C necessary for securing BH content. In this case, it is necessary to satisfy the condition of C- (12/48) Ti (*) ≦ 0.05%, preferably ≦ 0.03%.
Ｎｂは、Ｔｉ同様の効果がある。ただし、０．０１％未満ではこの効果が不十分であり、０．２％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、０．２％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過剰な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、さらにＢＨ量確保に必要な最適な固溶Ｃ量を得るためには、Ｔｉ（＊）＝Ｔｉ＋（４８／９３）Ｎｂ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％、かつＣ−（１２／４８）Ｔｉ（＊）≦０．０５％、好ましくは≦０．０３％の条件を満たすことが必要である。 Nb has the same effect as Ti. However, if the content is less than 0.01%, this effect is insufficient. Even if the content exceeds 0.2%, the effect is not only saturated but also the alloy cost is increased. Therefore, the Nb content is 0.01% or more and 0.2% or less. Furthermore, in order to contribute to the improvement of burring workability by precipitating and fixing excess C which causes carbides such as cementite which deteriorates burring workability, in order to obtain the optimum solute C content necessary for securing the BH content. Ti (*) = Ti + (48/93) Nb− (48/14) N− (48/32) S ≧ 0% and C− (12/48) Ti (*) ≦ 0.05% Preferably, it is necessary to satisfy the condition of ≦ 0.03%.
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 precipitates with Ti and Nb at higher temperatures than C, reducing the Ti and Nb effective to fix the desired C. Therefore, it should be reduced as much as possible, but 0.005% or less is an acceptable range.
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、０．０００５％未満添加してもその効果がなく、Ｃａならば０．００５％超、ＲＥＭならば０．０２％超添加してもその効果が飽和するので、Ｃａ：０．０００５〜０．００５％、ＲＥＭ：０．０００５〜０．０２％添加する。 Ca and REM are elements that are detoxified by changing the form of non-metallic inclusions that become the starting point of destruction or deteriorate workability. However, even if less than 0.0005% is added, there is no effect, and if Ca is more than 0.005%, and if REM is added more than 0.02%, the effect is saturated. Add 0.005%, REM: 0.0005-0.02%.
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、相対的に軟化する恐れがある場合に焼き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％未満ではそれら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 has an effect of increasing the fatigue limit by suppressing grain boundary embrittlement due to P, which is considered to be caused by a decrease in the amount of solute C. Therefore, B is added as necessary. Furthermore, when the base metal strength is 640 MPa or more, in the portion subjected to the thermal history in which α → γ → α transformation occurs in the weld heat affected zone, there is a risk that the material will not soften due to low Ceq and may be relatively softened. By adding B that improves the hardenability, softening at the relevant part is suppressed, and there is an effect of transitioning the fracture form of the joint from the welded part to the base material part. Therefore, it is added as necessary. However, if it is less than 0.0002%, it is insufficient for obtaining these effects, and if added over 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Ｍｏ，Ｖ，Ｃｒ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０．０５％、０．０２％、０．０１％未満ではその効果を得ることができない。また、それぞれ１％、０．２％、１％を超え添加してもその効果は飽和する。 Furthermore, in order to provide strength, one or more of precipitation strengthening or solid solution strengthening elements of Mo, V, and Cr may be added. However, the effect cannot be obtained if the content is less than 0.05%, 0.02%, and 0.01%, respectively. Moreover, the effect is saturated even if added over 1%, 0.2%, and 1%, respectively.
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可能性があるため、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 may occur during winding, and the static strength of the steel sheet may increase significantly due to precipitation strengthening. Will deteriorate significantly. Further, with such Cu precipitation strengthening, the fatigue limit ratio does not improve as much as the increase in static strength, so the fatigue limit ratio decreases. Therefore, the Cu content is 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、０．６％を超えて添加してもその効果が飽和するので、０．１〜０．６％とする。 Ni is added as necessary to prevent hot brittleness due to Cu inclusion. However, if less than 0.1%, the effect is small, and even if added over 0.6%, the effect is saturated, so 0.1% to 0.6%.
なお、これらを主成分とする鋼にＺｒ，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。 Note that Zr, Sn, Co, Zn, W, and Mg may be contained in a total amount of 1% or less in steel containing these as main components. However, since Sn may cause wrinkles during hot rolling, 0.05% or less is desirable.
次に本発明における鋼板のミクロ組織ついて詳細に説明する。
ＢＨ性とバーリング加工性とを両立させるためには、そのミクロ組織がベイニティックフェライトであり、旧オーステナイト粒の平均粒径が８μｍ超〜８０μｍであることが必要である。ここで、本発明おけるベイニティックフェライトは少量のγr 、ＭＡはその合計量を３％以下で含むことは許容される。優れたＢＨ性とバーリング加工性とを両立させるためには、上述したようにベイニティックフェライトが優れているが、鋼板のミクロ組織としてベイニティックフェライト以外にポリゴナルフェライトを少量ならば含んでもその特性を大幅に劣化させるものではないが、ＢＨ性を劣化させないためには最大２０％以下とすることが望ましい。
Next, the microstructure of the steel sheet in the present invention will be described in detail.
In order to achieve both BH properties and burring workability, the microstructure needs to be bainitic ferrite and the average grain size of prior austenite grains needs to be more than 8 μm to 80 μm. Here, the bainitic ferrite in the present invention is allowed to contain a small amount of γ r and MA in a total amount of 3% or less. In order to achieve both excellent BH properties and burring workability, bainitic ferrite is excellent as described above, but even if it contains a small amount of polygonal ferrite in addition to bainitic ferrite as a microstructure of the steel sheet. Although it does not significantly deteriorate the characteristics, it is desirable to make it 20% or less at maximum in order not to deteriorate the BH property.
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。
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、あるいは熱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。
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、転炉や電炉等による溶製に引き続き、各種の２次精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。
Next, the reasons for limiting the production method of the present invention will be described in detail below.
The present invention can also be obtained by casting, hot rolling after cooling or after hot rolling, or by subjecting hot-rolled steel sheets to heat treatment in a hot dipping line and further subjecting these steel sheets to surface treatment. It is done.
In the present invention, the production method preceding hot rolling is not particularly limited. That is, following smelting by blast furnace, converter, electric furnace, etc., the components are adjusted so that the desired component content is obtained by various secondary refining, and then, in addition to normal continuous casting, casting by ingot method, thin slab What is necessary is just to cast by methods, such as casting. Scrap may be used as a raw material. In the case of a slab obtained by continuous casting, it may be directly sent to a hot rolling mill as it is a high-temperature slab, or may be hot-rolled after being reheated in a heating furnace after being cooled to room temperature.
スラブ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℃未満の加熱ではスケジュール上操業効率を著しく損なうため、スラブ再加熱温度は１０００℃以上が望ましい。さらには、１１００℃未満の加熱ではスケールオフ量が少なくスラブ表層の介在物をスケールと共に後のデスケーリングによって除去できなくなる可能性があるので、スラブ再加熱温度は１１００℃以上が望ましい。 Although there is no restriction | limiting in particular about slab reheating temperature, Since a scale-off amount will become large and a yield will fall when it is 1400 degreeC or more, reheating temperature is desirable below 1400 degreeC. In addition, heating below 1000 ° C significantly impairs the operation efficiency in terms of schedule, so the slab reheating temperature is desirably 1000 ° C or higher. Furthermore, when the heating is less than 1100 ° C., the amount of scale-off is so small that inclusions on the surface of the slab cannot be removed together with the scale by subsequent descaling. Therefore, the slab reheating temperature is preferably 1100 ° C. or more.
熱間圧延工程は、粗圧延を終了後、仕上げ圧延を行うが、板厚方向によりベイニティックフェライトを得るためには仕上げ圧延開始温度を１０００℃以上が望ましい。さらにエッジ部の温度低下による幅方向の材質劣化を回避するためには、１０５０℃以上が望ましい。そのためには必要に応じて粗圧延終了から仕上圧延開始までの間または／および仕上圧延中に粗バーまたは圧延材を加熱する。 In the hot rolling process, finish rolling is performed after finishing rough rolling. In order to obtain bainitic ferrite in the sheet thickness direction, the finish rolling start temperature is desirably 1000 ° C. or higher. Furthermore, in order to avoid the material deterioration in the width direction due to the temperature drop of the edge portion, it is desirable that the temperature is 1050 ° C. or higher. For this purpose, the rough bar or the rolled material is heated as necessary from the end of rough rolling to the start of finish rolling or / and during finish rolling.
特に本発明のうちでも優れた破断延びを安定して得るためには、ＭｎＳ等の微細析出を抑制することが有効である。この場合の加熱装置はどのような方式でも構わないが、トランスバース型誘導加熱であれば板厚方向に均熱できるので、トランスバース型誘導加熱が望ましい。通常、ＭｎＳ等の析出物は１２５０℃程度のスラブ再加熱で再固溶が起こり、後の熱間圧延中に微細析出する。従って、スラブ再加熱温度を１１５０℃程度に制御し、ＭｎＳ等の再固溶を抑制できれば延性を改善できる。ただし、圧延終了温度を本発明の範囲にするためには、粗圧延終了から仕上圧延開始までの間または／および仕上げ圧延中での粗バーまたは圧延材の加熱が有効な手段となる。 In particular, in order to stably obtain excellent elongation at break in the present invention, it is effective to suppress fine precipitation of MnS or the like. Any heating apparatus may be used in this case, but transverse induction heating is desirable because transverse induction heating can equalize the thickness in the thickness direction. Usually, precipitates such as MnS are re-dissolved by slab reheating at about 1250 ° C. and finely precipitated during subsequent hot rolling. Therefore, ductility can be improved if the slab reheating temperature is controlled to about 1150 ° C. and re-solution of MnS or the like can be suppressed. However, in order to set the rolling end temperature within the range of the present invention, heating of the rough bar or rolled material from the end of rough rolling to the start of finish rolling and / or during finish rolling is an effective means.
粗圧延終了と仕上げ圧延開始の間にデスケーリングを行う場合は、鋼板表面での高圧水の衝突圧Ｐ（ＭＰａ）×流量Ｌ（リットル／ｃｍ2 ）≧０．００２５の条件を満たすことが望ましい。
鋼板表面での高圧水の衝突圧Ｐは以下のように記述される（「鉄と鋼」１９９１、ｖｏｌ．７７、Ｎｏ．９、ｐ１４５０参照）
Ｐ（ＭＰａ）＝５．６４×ＰO ×Ｖ／Ｈ2
ただし、
ＰO （ＭＰａ）：液圧力
Ｖ（リットル／ｍｉｎ）：ノズル流液量
Ｈ（ｃｍ）：鋼板表面とノズル間の距離
When descaling is performed between the end of rough rolling and the start of finish rolling, it is desirable to satisfy the condition of high-pressure water collision pressure P (MPa) × flow rate L (liters / cm 2 ) ≧ 0.0025 on the steel sheet surface. .
The collision pressure P of high-pressure water on the steel sheet surface is described as follows (see “Iron and Steel” 1991, vol. 77, No. 9, p1450).
P (MPa) = 5.64 × P O × V / H 2
However,
P O (MPa): Liquid pressure V (L / min): Nozzle flow rate H (cm): Distance between steel plate surface and nozzle
流量Ｌは以下のように記述される。
Ｌ（リットル／ｃｍ2 ）＝Ｖ／（Ｗ×ｖ）
ただし、
Ｖ（リットル／ｍｉｎ）：ノズル流液量
Ｗ（ｃｍ）：ノズル当たり噴射液が鋼板表面に当たっている幅
ｖ（ｃｍ／ｍｉｎ）：通板速度
衝突圧Ｐ×流量Ｌ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、ノズル流液量を増加させるとノズルの摩耗が激しくなる等の不都合が生じるため、０．０２以下とすることが望ましい。
The flow rate L is described as follows.
L (liter / cm 2 ) = V / (W × v)
However,
V (liter / min): Nozzle flow rate W (cm): Width of spray liquid per nozzle hitting steel plate surface v (cm / min): Plate passing speed The upper limit of collision pressure P × flow rate L is the effect of the present invention. Although it is not necessary to determine in particular in order to obtain it, it is desirable to make it 0.02 or less because an increase in the amount of nozzle flow causes problems such as increased wear on the nozzle.
さらに、仕上げ圧延後の鋼板表面の最大高さＲｙが１５μｍ（１５μｍＲｙ，ｌ２．５ｍｍ，ｌｎ１２．５ｍｍ）以下であることが望ましい。これは、例えば「金属材料疲労設計便覧」、日本材料学会編、８４頁に記載されている通り、熱延または酸洗ままの鋼板の疲労強度は鋼板表面の最大高さＲｙと相関があることから明らかである。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、５秒以内に行うのが望ましい。
また、粗圧延と仕上げ圧延の間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。
Furthermore, it is desirable that the maximum height Ry of the steel sheet surface after finish rolling is 15 μm (15 μm Ry, l2.5 mm, ln12.5 mm) or less. For example, as described in “Handbook of Fatigue Design for Metallic Materials”, edited by the Japan Society of Materials Science, page 84, the fatigue strength of a hot-rolled or pickled steel sheet has a correlation with the maximum height Ry of the steel sheet surface. It is clear from Further, the subsequent finish rolling is desirably performed within 5 seconds in order to prevent the scale from being generated again after descaling.
Moreover, a sheet bar may be joined between rough rolling and finish rolling, and finish rolling may be performed continuously. At that time, the coarse bar may be wound once in a coil shape, stored in a cover having a heat retaining function as necessary, and rewound again before joining.
仕上げ圧延は、当該成分系にて望ましいベイニティックフェライト組織を得るためには、圧延終了後のフェライトの析出を抑制する必要があるので、最終段での圧下率の合計が１〜１５％の圧延を行う必要がある。最終段の圧下率が１％未満では、鋼板の平坦度が劣化し、１５％超ではフェライトの析出が進行しすぎて、望ましいベイニティックフェライト組織を得られない。従って最終段の圧下率は１〜１５％とする。 In the finish rolling, in order to obtain the desired bainitic ferrite structure in the component system, it is necessary to suppress the precipitation of ferrite after the rolling, so the total rolling reduction in the final stage is 1 to 15%. It is necessary to perform rolling. If the rolling reduction at the final stage is less than 1%, the flatness of the steel sheet deteriorates, and if it exceeds 15%, precipitation of ferrite proceeds excessively, and a desirable bainitic ferrite structure cannot be obtained. Therefore, the rolling reduction in the final stage is 1 to 15%.
仕上げ圧延終了温度（ＦＴ）をＡｒ3 変態点温度＋５０℃以上とする。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×％Ｍｎｅｑ
ただし、Ｍｎｅｑ＝Ｍｎ＋Ｃｒ＋Ｃｕ＋Ｍｏ＋Ｎｉ／２＋１０（Ｎｂ−０．０２）
または、Ｍｎｅｑ＝Ｍｎ＋Ｃｒ＋Ｃｕ＋Ｍｏ＋Ｎｉ／２＋１０（Ｎｂ−０．０２）＋１：Ｂ添加の場合
Finish rolling end temperature (FT) is Ar3 transformation point temperature + 50 ° C. or higher. Here, the Ar3 transformation point temperature is simply shown in relation to the steel components by the following calculation formula, for example. That is,
However, Mneq = Mn + Cr + Cu + Mo + Ni / 2 + 10 (Nb−0.02)
Or, Mneq = Mn + Cr + Cu + Mo + Ni / 2 + 10 (Nb−0.02) +1: When B is added
仕上げ圧延終了温度（ＦＴ）は、Ａｒ3 変態点温度＋５０℃未満であるとフェライトの析出が進行し易くなり、目的とするベイニティックフェライト組織が得られなくなるので、Ａｒ3 変態点温度＋５０℃以上とする。仕上げ圧延終了温度（ＦＴ）の上限は特に設けないが、Ａｒ3 変態点温度＋２００℃超を得るためには、加熱炉温度または粗圧延終了から仕上圧延開始までの間または／および仕上げ圧延中での粗バーまたは圧延材の加熱が設備的に負荷が大きいので、その上限はＡｒ3 変態点温度＋２００℃以下が望ましい。 If the finish rolling finish temperature (FT) is less than Ar3 transformation point temperature + 50 ° C., precipitation of ferrite tends to proceed and the intended bainitic ferrite structure cannot be obtained. Therefore, Ar3 transformation point temperature + 50 ° C. or more. To do. There is no particular upper limit for the finish rolling finish temperature (FT), but in order to obtain Ar3 transformation temperature + 200 ° C., the heating furnace temperature or from the end of rough rolling to the start of finish rolling and / or during finish rolling Since the heating of the rough bar or the rolled material has a large load in terms of equipment, the upper limit is desirably Ar3 transformation point temperature + 200 ° C. or less.
仕上げ圧延終了後、Ａｒ3 〜Ｂｆの温度域を８０〜５００℃／sec 以上の冷却速度で冷却するが、Ａｒ3 変態点温度以上より冷却を開始しないとフェライトの析出が進行し、目的とするベイニティックフェライト組織が得られなくなる。従って、冷却はＡｒ3 変態点以上にて開始する。一方、Ｂｆ以上で冷却を停止すると、やはりベイニティックフェライトへの変態が不十分で、十分なベイニティックフェライト組織が得られなくなる恐れがある。従って、冷却する温度域はＡｒ3 〜Ｂｆである。 After finishing rolling, the temperature range of Ar3 to Bf is cooled at a cooling rate of 80 to 500 [deg.] C./sec or more, but if the cooling is not started above the Ar3 transformation point temperature, precipitation of ferrite proceeds and the target baini A tick ferrite structure cannot be obtained. Therefore, cooling starts above the Ar3 transformation point. On the other hand, if the cooling is stopped at Bf or more, the transformation to bainitic ferrite is still insufficient, and there is a possibility that a sufficient bainitic ferrite structure cannot be obtained. Therefore, the cooling temperature range is Ar3 to Bf.
ただし、仕上げ圧延終了後０．５〜５秒間で冷却を開始する。０．５秒未満に冷却を開始すると、オーステナイトの再結晶および粒成長が不十分となり、８μｍ超のオーステナイト粒径が得られなくなる恐れがあり、５秒超後に冷却を開始するとＡｒ3 変態点よりも温度が低下してしまう可能性があり、Ａｒ3 変態点以上で冷却を開始することができない恐れがあるので５秒以内とする。さらにオーステナイト粒成長を抑制し、粒径を８０μｍ以下にすると言う観点からは４秒以内が望ましい。
However, cooling is started 0.5 to 5 seconds after finish rolling. If cooling is started in less than 0.5 seconds, austenite recrystallization and grain growth may be insufficient, and an austenite grain size of more than 8 μm may not be obtained. If cooling is started after more than 5 seconds, the Ar3 transformation point may be exceeded. The temperature may drop, and cooling may not be started above the Ar3 transformation point, so the time is within 5 seconds. Further, from the viewpoint of suppressing the austenite grain growth and making the
また、冷却速度は８０℃／sec 未満では、フェライトの析出が進行し目的とするベイニティックフェライト組織が得られず、ＢＨ性が十分確保できない。従って、冷却速度は８０℃／sec 以上とする。さらに確実にベイニティックフェライト組織を得るためには１３０℃／sec 以上の冷却速度が望ましい。また、冷却速度が５００℃／sec 超ではマルテンサイト変態が進行し、目的とするミクロ組織得られないので、その上限は５００℃／sec 以下とする。さらに、熱ひずみによる板そりが懸念されることから、２５０℃／sec 以下とすることが望ましい。
ここでＢｆ温度（ベイニティックフェライトの変態終了温度）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
Ｂｆ＝７１０−２７０×％Ｃ−９０×％Ｍｎ−３７×％Ｎｉ−７０×％Ｃｕ
On the other hand, if the cooling rate is less than 80 ° C./sec, the precipitation of ferrite proceeds and the intended bainitic ferrite structure cannot be obtained, and the BH property cannot be sufficiently secured. Accordingly, the cooling rate is 80 ° C./sec or more. In order to obtain a bainitic ferrite structure more reliably, a cooling rate of 130 ° C./sec or more is desirable. On the other hand, if the cooling rate exceeds 500 ° C./sec, martensitic transformation proceeds and the desired microstructure cannot be obtained, so the upper limit is made 500 ° C./sec or less. Furthermore, since there is a concern about plate warpage due to thermal strain, it is desirable to set it to 250 ° C./sec or less.
Here, the Bf temperature (the transformation end temperature of bainitic ferrite) is simply shown in relation to the steel component by the following calculation formula, for example. That is, Bf = 710-270 ×% C-90 ×% Mn-37 ×% Ni-70 ×% Cu
上記温度域を特定の冷却速度で冷却した後に、２０℃／sec 以上の冷却速度で冷却した後に巻取る。この時の冷却速度が２０℃／sec 未満であると、十分な固溶Ｃが凍結できずＢＨ量が減少する。また上限は特に限定しないが、１００℃／sec 超の冷却速度で冷却しても効果が飽和するので、その上限は１００℃／sec 以下が望ましい。さらにＴｉ，Ｎｂ等による析出強化を有効に活用するために、巻取る前にＢｆ〜５００℃の温度域で２〜１２秒の空冷を行った後に当該冷却速度での冷却を行ってもよい。この温度が５００℃未満であると析出が十分に進行しない。また２秒未満ではやはり析出が不十分であり、１２秒超というのは通板速度、ランナウトテーブルの長さより非現実的である。さらに、ＢＨ量を十分に確保すると言う観点からは８秒以内が望ましい。 The temperature range is cooled at a specific cooling rate, and then cooled at a cooling rate of 20 ° C./sec or more, and then wound. If the cooling rate at this time is less than 20 ° C./sec, sufficient solid solution C cannot be frozen and the amount of BH decreases. Although the upper limit is not particularly limited, the effect is saturated even when cooled at a cooling rate exceeding 100 ° C./sec. Therefore, the upper limit is preferably 100 ° C./sec or less. Further, in order to effectively utilize precipitation strengthening by Ti, Nb, etc., cooling at the cooling rate may be performed after air cooling for 2 to 12 seconds in a temperature range of Bf to 500 ° C. before winding. If this temperature is less than 500 ° C., precipitation does not proceed sufficiently. Also, if it is less than 2 seconds, precipitation is still insufficient, and exceeding 12 seconds is more impractical than the plate feed speed and the runout table length. Further, from the viewpoint of securing a sufficient amount of BH, it is preferably within 8 seconds.
巻取温度は５００℃超では、当該温度域ではＣの拡散が容易であり、ＢＨ性を高める固溶ＣがＴｉ等に固定され十分確保できないため、巻取温度は５００℃以下に限定する。ＢＨ量を６０ＭＰａ以上得たい場合は３５０℃以下が望ましい。さらに望ましくは２００℃以下である。巻取温度の下限値は特に限定しないが、コイルが長時間水濡れの状態にあると錆による外観不良が懸念されるため、５０℃以上が望ましい。 When the coiling temperature is higher than 500 ° C., the diffusion of C is easy in the temperature range, and the solid solution C that improves the BH property is fixed to Ti or the like and cannot be sufficiently secured. When it is desired to obtain a BH amount of 60 MPa or more, 350 ° C. or less is desirable. More desirably, it is 200 ° C. or lower. The lower limit value of the coiling temperature is not particularly limited. However, if the coil is in a wet state for a long time, there is a concern about poor appearance due to rust.
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。 なお、鋼板形状の矯正や可動転位導入による延性の向上のためには、０．１％以上２％以下のスキンパス圧延を施すことが望ましい。
酸洗後の熱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漬し、必要に応じて合金化処理してもよい。
After completion of the hot rolling process, pickling may be performed as necessary, and then a skin pass with a reduction rate of 10% or less or cold rolling to a reduction rate of about 40% may be performed inline or offline. In order to improve the ductility by correcting the shape of the steel plate or introducing movable dislocations, it is desirable to perform skin pass rolling of 0.1% or more and 2% or less.
In order to galvanize the hot-rolled steel sheet after pickling, it may be immersed in a galvanizing bath and alloyed as necessary.
以下に、実施例により本発明をさらに説明する。
表２に示す化学成分を有するＡ〜Ｋの鋼は、転炉にて溶製して、連続鋳造後、直送もしくは再加熱し、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。また、鋼Ｇについては粗圧延後に衝突圧２．７ＭＰａ、流量０．００１リットル／ｃｍ2 の条件でデスケーリングを施した。さらに、表３に示すように鋼Ｂについては亜鉛めっきを施した。
The following examples further illustrate the present invention.
A to K steels having the chemical components shown in Table 2 are melted in a converter, continuously cast, then directly sent or reheated, and 1.2 to 5.5 mm in finish rolling following rough rolling. It was wound up after thickening. However, the display about the chemical composition in a table | surface is the mass%. Steel G was subjected to descaling after rough rolling under conditions of a collision pressure of 2.7 MPa and a flow rate of 0.001 liter / cm 2 . Furthermore, as shown in Table 3, the steel B was galvanized.
製造条件の詳細を表３に示す。ここで、「粗バー加熱」は粗圧延終了から仕上圧延開始までの間または／および仕上げ圧延中に粗バーまたは圧延材を加熱の有無を、「ＦＴ０」は仕上げ圧延温度開始、「ＦＴ」は仕上げ圧延温度終了、「冷却開始までの時間」とは仕上げ圧延終了から冷却を開始するまでの時間を、「Ａｒ3 〜Ｂｆでの冷却速度」とは冷却時にＡｒ3 〜Ｂｆの温度域を通過する時の平均冷却速度を、「Ｂｆ未満での冷却速度」とは冷却時にＢｆ未満の温度域を通過する時の平均冷却速度を、「ＣＴ」とは巻取温度を示している。 Details of the manufacturing conditions are shown in Table 3. Here, “rough bar heating” means whether or not the rough bar or rolled material is heated during the period from the end of rough rolling to the start of finish rolling or / and during finish rolling, “FT0” is the start of finish rolling temperature, and “FT” is End of finish rolling temperature, "Time to start cooling" means time from finish finish rolling to start of cooling, "Cooling rate at Ar3 to Bf" means to pass through the temperature range of Ar3 to Bf during cooling "Cooling rate below Bf" means the average cooling rate when passing through a temperature range below Bf during cooling, and "CT" shows the coiling temperature.
このようにして得られた薄鋼板の引張試験は、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。
ＢＨ試験は引張試験と同様にＪＩＳ Ｚ ２２０１に記載の５号試験片に加工し、２％の引張予ひずみを試験片に付与した後、１７０℃×２０分の塗装焼き付け工程相当の熱処理を施してから再度引張試験を実施した。ここでＢＨ量とは、再引張での上降伏点から２％の引張り予ひずみの流動応力を差し引いたと定義される。
バーリング加工性は、日本鉄鋼連盟規格ＪＦＳ Ｔ １００１−１９９６記載の穴拡げ試験方法に従い、穴拡げ値にて評価した。
The tensile test of the thin steel plate thus obtained was performed by first processing the specimen into a No. 5 test piece described in JIS Z 2201, and following the test method described in JIS Z 2241.
The BH test is processed into a No. 5 test piece described in JIS Z 2201 in the same way as the tensile test, 2% tensile pre-strain is applied to the test piece, and then a heat treatment equivalent to a coating baking process of 170 ° C. × 20 minutes is performed. Then, the tensile test was performed again. Here, the BH amount is defined as subtracting 2% tensile prestrained flow stress from the upper yield point in re-tensioning.
Burring workability was evaluated by a hole expansion value according to a hole expansion test method described in Japan Iron and Steel Federation Standard JFS T 1001-1996.
一方、ミクロ組織の調査は鋼板板幅の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表層下０．２ｍｍ、板厚の１／４ｔ、１／２ｔにおける視野の写真にて行った。ミクロ組織の体積分率とは上記金属組織写真において面積分率で定義される。 On the other hand, the microstructure was examined by grinding a sample cut from a 1/4 W or 3/4 W position of the steel plate width to a cross section in the rolling direction, etching using a Nital reagent, and 200-500 times magnification using an optical microscope. The observation was carried out with photographs of the field of view at 0.2 mm below the surface layer and at 1/4 t and 1/2 t of the plate thickness. The volume fraction of the microstructure is defined by the area fraction in the metal structure photograph.
ここでベイニティックフェライトとは、日本鉄鋼協会基礎研究会ベイナイト調査研究部会編、「低炭素鋼のベイナイト組織と変態挙動に関する最近の研究」、ベイナイト調査研究部会最終報告書（１９９４年、日本鉄鋼協会）に記載されているように、無拡散でせん断的機構により生成する変態組織と定義されるミクロ組織である。 Here, bainitic ferrite refers to the latest research on bainite structure and transformation behavior of low-carbon steel, edited by the Japan Iron and Steel Institute Basic Research Group, Bainite Research Group, 1994 Bainite Research Group Final Report (1994, Japan Steel). This is a microstructure defined as a metamorphic structure generated by a non-diffusion and shearing mechanism, as described in the Association).
次に旧オーステナイト粒の平均粒径の測定であるが、ナイタール試薬にてエッチングし光学顕微鏡にて観察した同一試料を再度研磨し、特開平０６−２０７２７９に記載の腐食液および腐食方法に従いエッチングした後、ＪＩＳ Ｇ ０５５２記載の切断法を用い、その測定値より求めた粒度番号Ｇより、断面積１ｍｍ2 当たりの結晶粒の数ｍをｍ＝８×２G より求め、このｍよりｄm ＝１／√ｍで得られる平均粒径ｄm を平均粒径と定義する。 Next, the average particle size of the prior austenite grains was measured. The same sample that was etched with a Nital reagent and observed with an optical microscope was polished again, and etched according to the corrosive solution and the corrosive method described in JP-A-06-207279. Thereafter, using the cutting method described in JIS G 0552, the number m of crystal grains per 1 mm 2 in cross-sectional area is obtained from m = 8 × 2 G from the particle size number G obtained from the measured value. From this m, d m = the average particle size d m obtained by 1 / √m defined as the average particle size.
本発明に沿うものは、鋼Ａ−１、Ａ−２、Ｂ、Ｃ、Ｄ、Ｅ、Ｆ、Ｇの８鋼であり、所定の量の鋼成分を含有し、そのミクロ組織が主にベイニティックフェライトであり、旧オーステナイト粒の平均粒径が８μｍ超〜８０μｍであることを特徴とする、焼付け硬化性を有する高バーリング熱延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穴拡げ値およびＢＨ量がそれぞれ７０％、５０ＭＰａを上回っている。 In accordance with the present invention, steels A-1, A-2, B, C, D, E, F, and G are 8 steels, containing a predetermined amount of steel components, and the microstructure is mainly bay. A high burring hot-rolled steel sheet having bake hardenability is obtained, characterized in that it is nittic ferrite and the average grain size of prior austenite grains is more than 8 μm to 80 μm. The evaluated hole expansion value and BH amount exceed 70% and 50 MPa, respectively.
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Ａ−３は、仕上げ圧延終了温度（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ＢＨ量が得られていない。鋼Ａ−４は、最終段での圧下率が本発明請求項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ＢＨ量が得られていない。鋼Ａ−５は、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ＢＨ量が得られていない。鋼Ａ−６は、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値が得られていない。
Steels other than the above are outside the scope of the present invention for the following reasons.
That is, since the finish rolling finish temperature (FT) of steel A-3 is outside the range of claim 8 of the present invention, the target microstructure of
鋼Ａ−７は、Ａｒ3 〜Ｂｆでの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ＢＨ量が得られていない。鋼Ａ−８は、Ｂｆ未満からの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値とＢＨ量が得られていない。鋼Ａ−９は、巻取温度（ＣＴ）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ＢＨ量が得られていない。鋼Ｈは、Ｃの含有量が本発明請求項１の範囲外であり、Ａｒ3 〜Ｂｆでの冷却速度と巻取温度（ＣＴ）も本発明請求項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、十分な強度とＢＨ量が得られていない。
In Steel A-7, since the cooling rate at Ar3 to Bf is outside the range of Claim 8 of the present invention, the target microstructure of
鋼Ｉは、Ｃ−（１２／４８）Ｔｉ（＊）の値が本発明請求項１の範囲外であり、Ａｒ3 〜Ｂｆでの冷却速度と巻取温度（ＣＴ）も本発明請求項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、十分な穴拡げ値とＢＨ量が得られていない。鋼Ｊは、Ｔｉの含有量が本発明請求項１の範囲外であり、かつ、Ａｒ3 〜Ｂｆでの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、十分な穴拡げ値とＢＨ量が得られていない。鋼Ｋは、Ｃの含有量が本発明請求項１の範囲外でかつＴｉが無添加あり、Ａｒ3 〜Ｂｆでの冷却速度も本発明請求項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、十分な穴拡げ値とＢＨ量が得られていない。
Steel I has a value of C- (12/48) Ti (*) outside the range of
Claims (13)
Ｃ ：０．０１〜０．１％、
Ｓｉ：０．００５〜２％、
Ｍｎ：０．１〜３％、
Ｐ ≦０．１％、
Ｓ ≦０．０３％、
Ａｌ：０．００１〜０．１％、
Ｎ ≦０．００５％、
Ｔｉ：０．０５〜０．２％
を含み、さらに
Ｔｉ（＊）＝Ｔｉ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％
かつ
Ｃ−（１２／４８）Ｔｉ（＊）≦０．０５％
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板であって、そのミクロ組織が主にベイニティックフェライトであり、旧オーステナイト粒の平均粒径が８μｍ超〜８０μｍであることを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。 In mass%
C: 0.01 to 0.1%,
Si: 0.005 to 2%,
Mn: 0.1 to 3%
P ≦ 0.1%,
S ≦ 0.03%,
Al: 0.001 to 0.1%,
N ≦ 0.005%,
Ti: 0.05 to 0.2%
Ti (*) = Ti− (48/14) N− (48/32) S ≧ 0%
And C- (12/48) Ti (*) ≦ 0.05%
Is a steel plate that contains C, S, N, Ti in the range satisfying the following, the balance is Fe and inevitable impurities, the microstructure is mainly bainitic ferrite, the average grain size of the prior austenite grains is A high burring hot-rolled steel sheet having bake hardenability, characterized by being over 8 μm to 80 μm.
Ｎｂ：０．０１〜０．２％
を含み、さらに
Ｔｉ（＊）＝Ｔｉ＋（４８／９３）Ｎｂ−（４８／１４）Ｎ−（４８／３２）Ｓ≧０％かつ
Ｃ−（１２／４８）Ｔｉ（＊）≦０．０５％
を満たす範囲でＮｂを含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。 The steel according to claim 1, further in mass%,
Nb: 0.01 to 0.2%
Ti (*) = Ti + (48/93) Nb− (48/14) N− (48/32) S ≧ 0% and C− (12/48) Ti (*) ≦ 0.05%
A high burring hot-rolled steel sheet having bake hardenability characterized by containing Nb in a range satisfying the above.
Ｃａ：０．０００５〜０．００５％、
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焼付け硬化性を有する高バーリング熱延鋼板。 The steel according to claim 1 or 2, further in mass%,
Ca: 0.0005 to 0.005%,
REM: 0.0005 to 0.02%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or two of the following.
Ｂ ：０．０００２〜０．００２％
を含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。 The steel according to any one of claims 1 to 3, further in mass%,
B: 0.0002 to 0.002%
A high burring hot-rolled steel sheet having bake hardenability, comprising:
Ｍｏ：０．０５〜１％、
Ｖ ：０．０２〜０．２％、
Ｃｒ：０．０１〜１％
の一種または二種以上を含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。 The steel according to any one of claims 1 to 4, further in mass%,
Mo: 0.05 to 1%
V: 0.02-0.2%,
Cr: 0.01 to 1%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or more of the above.
Ｃｕ：０．２〜１．２％、
Ｎｉ：０．１〜０．６％
の一種または二種を含有することを特徴とする焼付け硬化性を有する高バーリング熱延鋼板。 The steel according to any one of claims 1 to 5, further in mass%,
Cu: 0.2 to 1.2%,
Ni: 0.1 to 0.6%
A high burring hot-rolled steel sheet having bake hardenability, characterized by containing one or two of the following.
A method for producing a high burring hot-rolled steel sheet having bake hardenability, wherein the alloying treatment is performed after galvanization in the production method according to claim 12.
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