JP2008266695A - High strength hot-rolled steel sheet and producing method therefor - Google Patents

High strength hot-rolled steel sheet and producing method therefor Download PDF

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JP2008266695A
JP2008266695A JP2007108759A JP2007108759A JP2008266695A JP 2008266695 A JP2008266695 A JP 2008266695A JP 2007108759 A JP2007108759 A JP 2007108759A JP 2007108759 A JP2007108759 A JP 2007108759A JP 2008266695 A JP2008266695 A JP 2008266695A
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steel sheet
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rolled steel
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JP5214905B2 (en
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Kazuaki Hakomori
一昭 箱守
Yuji Kusumoto
雄二 楠本
Fuyuki Yoshida
冬樹 吉田
Ichiro Takeshi
伊知郎 竹士
Takashi Otani
崇 大谷
Takaaki Kurahashi
隆朗 倉橋
Masahiko Oda
昌彦 織田
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Nakayama Steel Works Ltd
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Priority to US12/078,860 priority patent/US20090223609A1/en
Priority to KR1020080033714A priority patent/KR101446354B1/en
Priority to EP08154615A priority patent/EP1990430B1/en
Priority to CN2008100922456A priority patent/CN101319295B/en
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    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/40—Ferrous alloys, e.g. steel alloys containing chromium with nickel
    • C22C38/44—Ferrous alloys, e.g. steel alloys containing chromium with nickel with molybdenum or tungsten
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/02—Ferrous alloys, e.g. steel alloys containing silicon
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/002—Heat treatment of ferrous alloys containing Cr
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D6/00—Heat treatment of ferrous alloys
    • C21D6/008—Heat treatment of ferrous alloys containing Si
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0221—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the working steps
    • C21D8/0226—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/0247—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment
    • C21D8/0263—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips characterised by the heat treatment following hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D8/00—Modifying the physical properties by deformation combined with, or followed by, heat treatment
    • C21D8/02—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips
    • C21D8/04—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing
    • C21D8/0421—Modifying the physical properties by deformation combined with, or followed by, heat treatment during manufacturing of plates or strips to produce plates or strips for deep-drawing characterised by the working steps
    • C21D8/0426—Hot rolling
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D9/00—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor
    • C21D9/46—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals
    • C21D9/48—Heat treatment, e.g. annealing, hardening, quenching or tempering, adapted for particular articles; Furnaces therefor for sheet metals deep-drawing sheets
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/001—Ferrous alloys, e.g. steel alloys containing N
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/04—Ferrous alloys, e.g. steel alloys containing manganese
    • C—CHEMISTRY; METALLURGY
    • C22—METALLURGY; FERROUS OR NON-FERROUS ALLOYS; TREATMENT OF ALLOYS OR NON-FERROUS METALS
    • C22C—ALLOYS
    • C22C38/00—Ferrous alloys, e.g. steel alloys
    • C22C38/18—Ferrous alloys, e.g. steel alloys containing chromium
    • C22C38/34—Ferrous alloys, e.g. steel alloys containing chromium with more than 1.5% by weight of silicon
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/001—Austenite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/002—Bainite
    • C—CHEMISTRY; METALLURGY
    • C21—METALLURGY OF IRON
    • C21D—MODIFYING THE PHYSICAL STRUCTURE OF FERROUS METALS; GENERAL DEVICES FOR HEAT TREATMENT OF FERROUS OR NON-FERROUS METALS OR ALLOYS; MAKING METAL MALLEABLE, e.g. BY DECARBURISATION OR TEMPERING
    • C21D2211/00—Microstructure comprising significant phases
    • C21D2211/008—Martensite

Abstract

<P>PROBLEM TO BE SOLVED: To provide a new high strength Si-Cr containing hot-rolled steel sheet having both high strength and good workability, and a producing method therefor. <P>SOLUTION: The high strength steel sheet having ≤1 μm the size and 5-20% volume ratio of uniformly dispersed retained austenite particles and mainly composed of bainite structure, is obtained by controlling so that Si and Cr are contained as the chemical component and old austenite particle diameter is ≤10 μm and by limiting the wounding temperature, wherein the volume ratio of martensite structure is ≤10%. <P>COPYRIGHT: (C)2009,JPO&INPIT

Description

請求項に係る発明は、高い引張強度をもちながらも優れた加工性を有する高強度熱延鋼板と、その製造方法に関するものである。   The invention according to the claims relates to a high-strength hot-rolled steel sheet having high workability while having high tensile strength, and a method for producing the same.

加工性の優れた高強度鋼板に対する最近の要請を、自動車の場合を例にして述べる。地球環境保全の観点から、自動車分野においてもＣＯ２等の排ガス量を低減していくことが是非とも必要である。そのためには、自動車車体の一層の軽量化が不可欠になる。車体の軽量化を達成するためには、自動車に使用される鋼板の強度を高めて、板厚を薄くしていかなければならない。同時に、自動車においては、搭乗者の安全性を確保していかなければならない。このためにも、鋼板の強度を一層高めていくことが必要になる。 The recent demand for high-strength steel sheets with excellent workability will be described by taking the case of automobiles as an example. From the viewpoint of protecting the global environment, it is necessary to reduce the amount of exhaust gas such as CO 2 in the automobile field. For that purpose, further weight reduction of the automobile body becomes indispensable. In order to reduce the weight of the car body, it is necessary to increase the strength of steel plates used in automobiles and reduce the thickness. At the same time, passengers must ensure the safety of passengers. For this purpose, it is necessary to further increase the strength of the steel sheet.

鋼板の強度が高くなると加工性が悪くなり、通常のプレス成形等の冷間加工法では高強度鋼板の適用が困難である。
ホットプレス法は、熱間でプレス加工をするのでスプリングバックの発生量は極めて少なく、形状凍結性が良い。そして、ホットプレスの際の焼入れ効果で、非常に高い強度をもった部品を高精度で提供することができる。しかしながら、ホットプレス加工前に鋼板を加熱することが必要であり、また、ホットプレス後にスケールを落とす作業が必要である。従って、作業効率が非常に悪い方法である。さらに、金型が加熱した鋼板と接するため金型の寿命が短いことも欠点であり、これが製造コストを増加させることにもなる。
ホットプレス後の鋼板は伸び値が小さく、部材が変形を受けた際に僅かな変形でも破断することがあるので、衝撃吸収能力が小さいと評価されている。従って、ホットプレス部品を、自動車等の重要保安部品として使用することは非常に難しい。 When the strength of the steel plate increases, the workability deteriorates, and it is difficult to apply a high strength steel plate by a cold working method such as ordinary press forming.
In the hot pressing method, since the hot pressing is performed, the amount of spring back is extremely small and the shape freezing property is good. And the part with very high intensity | strength can be provided with high precision by the quenching effect in the case of a hot press. However, it is necessary to heat the steel plate before hot pressing, and an operation of dropping the scale after hot pressing is required. Therefore, the work efficiency is very poor. Furthermore, since the mold contacts the heated steel plate, the short life of the mold is also a drawback, which increases the manufacturing cost.
The steel sheet after hot pressing has a small elongation value, and when the member is deformed, it may be broken even if it is slightly deformed. Therefore, it is very difficult to use a hot press part as an important safety part for an automobile or the like.

強度を高める方法としては、固溶強化、析出強化、結晶粒微細化強化および低温変態組織利用による強化などが基本的な方法である。固溶強化や析出強化といった多量の合金添加を必要とする強化機構の適用だけでは、極めて高い強度を必要とする鋼板の製造は不可能である。また結晶粒微細化強化を適用するにしても、強度の上昇はある程度図れても限界がある。低温変態組織利用による強化は１２００ＭＰａ超の鋼鈑を製造するには極めて有効な方法であるが、強度上昇に見合う延性の向上は期待できない。   Basic methods for increasing the strength include solid solution strengthening, precipitation strengthening, grain refinement strengthening, and strengthening by using a low temperature transformation structure. Only by applying a strengthening mechanism that requires a large amount of alloy addition such as solid solution strengthening or precipitation strengthening, it is impossible to produce a steel sheet that requires extremely high strength. Even if the grain refinement strengthening is applied, there is a limit even if the strength can be increased to some extent. Strengthening by using a low-temperature transformation structure is an extremely effective method for producing a steel plate of over 1200 MPa, but it cannot be expected to improve the ductility to meet the increase in strength.

一般的に、鋼板の強度を高めると、延性は小さくなり加工性は低くなる。
高強度鋼板の延性を高めた鋼板として、フェライトとマルテンサイト組織からなる複合組織（Ｄｕａｌ Ｐｈａｓｅ）鋼板、フェライト、ベイナイトと残留オーステナイト組織からなるＴＲＩＰ（Ｔｒａｎｓｆｏｒｍａｔｉｏｎ Ｉｎｄｕｃｅｄ Ｐｌａｓｔｉｃｉｔｙ）鋼板とよばれているものがある。
複合組織鋼板は、フェライト中に硬質なマルテンサイトを微細に分散させるが、この硬質なマルテンサイトにより、変形時に大きな加工硬化を引き起こし、高い延性を鋼板にもたらすのである。
ＴＲＩＰ鋼板については特許文献１、２にその例が示されている。残留オーステナイトを含有するこの種の鋼板は、その量と変形に対する安定度に応じて、加工誘起変態に起因する極めて良好な延性と成形性を有する。 Generally, when the strength of a steel plate is increased, ductility is reduced and workability is reduced.
Steel plates with high ductility of high-strength steel plates are called composite phase (Dual Phase) steel plates made of ferrite and martensite, and TRIP (Transformation Induced Plasticity) steel plates made of ferrite, bainite and retained austenite structures. .
In the composite steel sheet, hard martensite is finely dispersed in ferrite, and this hard martensite causes large work hardening at the time of deformation, and brings high ductility to the steel sheet.
Examples of the TRIP steel sheet are shown in Patent Documents 1 and 2. This type of steel sheet containing retained austenite has very good ductility and formability due to work-induced transformation, depending on its amount and stability to deformation.

さらに、鋼板の強度を１２００ＭＰａ以上に高めると、遅れ破壊の問題も発生する。遅れ破壊とは、部材の加工、組み立ての際には割れや破壊が発生せず、使用中に突如として割れが発生する現象のことである。特許文献３に示す高強度鋼板は、ベイナイトや焼戻しマルテンサイトなどの硬質な低温変態相に対し、フェライトのような軟質相を極力低減し、かつ残留オーステナイトの体積率を４％以下に制限することで、良好な耐遅れ破壊特性を確立したものである。
特開昭６０−４３４２５号公報 特開平９−１０４９４７号公報 特許第３２４７９０８号号公報 Furthermore, if the strength of the steel plate is increased to 1200 MPa or more, the problem of delayed fracture also occurs. Delayed fracture is a phenomenon in which cracks and breakage do not occur during the processing and assembly of members, but suddenly occur during use. The high-strength steel sheet shown in Patent Document 3 reduces the soft phase such as ferrite as much as possible to the hard low-temperature transformation phase such as bainite and tempered martensite, and limits the volume fraction of retained austenite to 4% or less. Thus, good delayed fracture resistance is established.
JP 60-43425 A JP-A-9-104947 Japanese Patent No. 3247908

高い強度を有しながら冷間加工での伸び特性を向上させた鋼板として、前記複合組織鋼板とＴＲＩＰ鋼板が挙げられる。   Examples of the steel sheet having high strength and improved elongation characteristics in cold working include the above-mentioned composite structure steel sheet and TRIP steel sheet.

複合組織鋼板では比較的低い合金添加量でも高い強度が得られ、同時に、加工硬化により良い均一伸び特性が得られる。
ＴＲＩＰ鋼板はさらに高い延性を示し、かつ高深絞り性を有するものである。そのため複雑な形状で高い加工性を必要とし、高い強度が要求される部材への適用が指向されている。
特許文献１のＴＲＩＰ鋼板は、圧延終了後の冷却工程で、４５０〜６５０℃の温度範囲で４〜２０秒保持し、オーステナイト中にフェライトを生成させた後、３５０℃以下まで冷却し、巻取る工程で製造する。
特許文献２では圧延終了後の冷却過程で、オーステナイト中にフェライトの生成を促進するため、Ａｒ３〜Ａｒ１での緩冷却を行うか、もしくは圧延完了温度をＡｒ３点近傍とし、その後３５０〜５００℃の範囲まで冷却し、巻取ることで製造する。
これらＴＲＩＰ鋼板はフェライト母相中にマルテンサイトもしくは残留オーステナイト、ベイナイトが分散した組織を有し、優れた強度と伸び特性を有する。
しかし、スポット溶接性が確保可能なＣ≦０．２０％では、引張強度で８００ＭＰａ程度しか得られず、加工性が渇望される、さらに高い強度範囲の鋼板の製造が困難である。
圧延終了後、途中に緩冷却を行わず、連続的に５００℃以下まで冷却する方法においても、圧延終了温度をＡｒ３点近傍とすれば、微細なフェライトの生成促進が可能となるが、Ａｒ３点近傍で圧延をした熱延鋼板の材質特性は、異方性が大きい問題がある。 In the composite structure steel plate, high strength can be obtained even with a relatively low alloy addition amount, and at the same time, good uniform elongation characteristics can be obtained by work hardening.
The TRIP steel sheet exhibits higher ductility and has high deep drawability. Therefore, application to a member that requires high workability with a complicated shape and requires high strength is directed.
The TRIP steel sheet of Patent Document 1 is held in a temperature range of 450 to 650 ° C. for 4 to 20 seconds in the cooling step after the rolling, and after generating ferrite in austenite, the steel is cooled to 350 ° C. or lower and wound. Produced in the process.
In Patent Document 2, in order to promote the formation of ferrite in the austenite in the cooling process after the end of rolling, slow cooling is performed with Ar3 to Ar1, or the rolling completion temperature is set to the vicinity of the Ar3 point, and then 350 to 500 ° C. Manufactured by cooling to range and winding.
These TRIP steel sheets have a structure in which martensite, retained austenite, and bainite are dispersed in a ferrite matrix, and have excellent strength and elongation characteristics.
However, if C ≦ 0.20%, where spot weldability can be ensured, only a tensile strength of about 800 MPa can be obtained, and it is difficult to produce a steel plate with a higher strength range that is craving for workability.
Even in the method of continuously cooling to 500 ° C. or less without performing slow cooling in the middle after the end of rolling, if the rolling end temperature is in the vicinity of the Ar3 point, the formation of fine ferrite can be promoted. The material property of the hot rolled steel sheet rolled in the vicinity has a problem of large anisotropy.

さらに、特許文献１に記載の熱延鋼板は、圧延加工度が低く、またＡ１点付近で冷却を一時停止するため、粗大なフェライト粒と残留オーステナイト粒が隣接した金属組織を示す。
遅れ破壊の原因である鋼板中に固溶した水素は、結晶構造に起因し、残留オーステナイト中に優先的にトラップされる。特に加工の影響を受け、加工誘起変態したマルテンサイトとフェライトの界面が最も危険なトラップサイトとされる。
残留オーステナイト粒が粗大であればあるほど、残留オーステナイト粒の体積に比べ、加工誘起変態したマルテンサイトとフェライトの界面の面積比が減少し、トラップされる水素濃度が高濃度化し、遅れ破壊の危険性が高まる。さらにマルテンサイトと残留オーステナイトが隣接した状態（ＭＡ）で共存していれば、破壊の伝播が促進され、さらに危険性が高まるとされる。
特許文献３に記載した高強度鋼板は、この残留オーステナイト量を制約することにより、耐遅れ破壊性を向上させたものである。しかし、高い強度を有しつつ、優れた加工性を得るためには、残留オーステナイトの活用は有効であり、その制約を設けずとも、遅れ破壊に対して無害化することが望ましい。 Furthermore, the hot-rolled steel sheet described in Patent Document 1 has a low degree of rolling and temporarily stops cooling in the vicinity of the point A1, and thus exhibits a metal structure in which coarse ferrite grains and residual austenite grains are adjacent.
Hydrogen dissolved in the steel sheet, which is the cause of delayed fracture, is preferentially trapped in the retained austenite due to the crystal structure. In particular, the interface between martensite and ferrite that has undergone processing-induced transformation is the most dangerous trap site due to the influence of processing.
The coarser the retained austenite grains, the smaller the ratio of the area of martensite-ferrite interface that has undergone work-induced transformation compared to the volume of retained austenite grains, the higher the concentration of trapped hydrogen, and the risk of delayed fracture Increases nature. Furthermore, if martensite and retained austenite coexist in a state of being adjacent (MA), propagation of fracture is promoted, and the risk is further increased.
The high-strength steel sheet described in Patent Document 3 has improved delayed fracture resistance by restricting the amount of retained austenite. However, in order to obtain excellent workability while having high strength, the utilization of retained austenite is effective, and it is desirable to make it harmless against delayed fracture without providing the restrictions.

そこで本願の発明者らは、粒径が１μｍ以下の残留オーステナイト（体積率５％以上２０％以下）を１０μｍ平方に７個以上微細分散させたベイナイト組織を得ることで、高い強度と良好な加工性及び耐遅れ破壊特性を併せもつ新しい低合金・高強度の鋼板およびその製法を開発したものである。   Therefore, the inventors of the present application obtain high strength and good processing by obtaining a bainite structure in which 7 or more of retained austenite having a particle size of 1 μm or less (volume ratio of 5% or more and 20% or less) is finely dispersed in 10 μm square. Has developed a new low-alloy / high-strength steel sheet that has both high reliability and delayed fracture resistance and its manufacturing method.

鋭意研究を行った結果、発明者らは、適正な圧延条件及び成分組成の採用で、好ましい高強度鋼板が得られることを見出した。すなわち、適正な成分範囲を有するスラブを、熱間圧延の粗圧延での高圧下圧延、仕上げ圧延での後段高ひずみ圧延を高温で終了し、数秒の空冷をした後に冷却を開始し、適正な温度で巻き取ることで、低合金組成で高い強度と優れた延性及び耐遅れ破壊特性を鋼板に同時に付与することが出来きるのである。その詳細を以下に示す。   As a result of intensive studies, the inventors have found that preferable high-strength steel sheets can be obtained by adopting appropriate rolling conditions and component compositions. That is, a slab having an appropriate component range is subjected to high pressure rolling in hot rolling rough rolling, subsequent high strain rolling in finish rolling at high temperature, air cooling for several seconds, and then cooling is started. By winding at temperature, it is possible to simultaneously impart high strength, excellent ductility and delayed fracture resistance to the steel sheet with a low alloy composition. Details are shown below.

請求項に記載した高強度鋼板は、残留オーステナイト粒の大きさが１μｍ以下でその体積率が５％以上２０％以下で、マルテンサイト組織の体積率が０％以上・１０％以下で、残部がベイナイト組織からなることを特徴とする高強度鋼板である。この時、旧オーステナイト粒径が１０μｍ以下で、その平均アスペクト比が２．０以下であれば、さらに望ましい特性を得ることが出来る。   The high-strength steel sheet described in the claims has a residual austenite grain size of 1 μm or less, a volume ratio of 5% to 20%, a martensite volume ratio of 0% to 10%, and the balance It is a high-strength steel plate characterized by comprising a bainite structure. At this time, if the prior austenite grain size is 10 μm or less and the average aspect ratio is 2.0 or less, more desirable characteristics can be obtained.

熱間圧延後のオーステナイト結晶粒を１０μｍ以下（図３）にすることによりベイナイト組織のラスを微細にし、さらに均一にベイナイト変態を完了させることにより、残留オーステナイトの大きさを１μｍ以下で効果的に組織内に１０μｍ平方に７個以上微細に分散させる。（図８）
このことにより多量の残留オーステナイトによる加工誘起塑性を利用した高延性鋼であるにもかかわらず、優れた耐遅れ破壊性を得ることが出来るのである。
また、旧オーステナイト粒のアスペクト比を２．０以下にする（図３）ことにより、圧延方向及び圧延直角方向に引張った材質の異方性が低減出来、さらなる加工性の向上が図れるのである。（図４） By reducing the austenite crystal grains after hot rolling to 10 μm or less (FIG. 3), the lath of the bainite structure is made finer, and by completing the bainite transformation uniformly, the size of the retained austenite is effectively reduced to 1 μm or less. 7 or more finely dispersed in 10 μm square in the tissue. (Fig. 8)
This makes it possible to obtain excellent delayed fracture resistance despite the fact that the steel is a high ductility steel utilizing work-induced plasticity caused by a large amount of retained austenite.
Further, by setting the aspect ratio of the prior austenite grains to 2.0 or less (FIG. 3), the anisotropy of the material pulled in the rolling direction and the direction perpendicular to the rolling can be reduced, and the workability can be further improved. (Fig. 4)

請求項に記載した高強度鋼板の成分範囲は、
Ｃ：０．１３〜０．２１、Ｓｉ：０．５〜２．０、Ｍｎ：０．２〜１．０
Ｃｒ：１.０〜４．０、Ｎｉ：０．０２〜１．０、 Ｍｏ：０.０５〜０．４
Ｐ：〜０．０１０、 Ｓ：〜０．００３、Ｎ：０．００５〜０．０１５
（各重量％）を含み、残部はＦｅおよび不可避的不純物の組成とするのがよい。
こうした適切な種類と量の化学成分を含むこととすれば、上記の組織を有していて望ましい機械的性質を発揮する高強度鋼板とすることが容易である。
合金元素は熱間圧延後の冷却およびその後の巻き取り過程で本発明の望ましい鋼板組織を得るため、ベイナイト変態に大きく影響するＣｒやＳｉを主要元素とする。これらの元素量を調節することでベイナイト変態を促進させ、マルテンサイトの形成量を抑制して、目的とする強度に制御することが可能なのである。
なお、各成分の作用については後述する。 The component range of the high-strength steel sheet described in the claim is:
C: 0.13-0.21, Si: 0.5-2.0, Mn: 0.2-1.0
Cr: 1.0-4.0, Ni: 0.02-1.0, Mo: 0.05-0.4
P: -0.010, S: -0.003, N: 0.005-0.015
(The respective weight percentages) are included, and the balance is preferably composed of Fe and inevitable impurities.
If such an appropriate kind and amount of chemical components are included, it is easy to obtain a high-strength steel sheet having the above-described structure and exhibiting desirable mechanical properties.
In order to obtain a desirable steel sheet structure of the present invention in the cooling process after hot rolling and the subsequent winding process, the alloying element is mainly composed of Cr or Si that greatly affects the bainite transformation. By adjusting the amounts of these elements, the bainite transformation is promoted, the amount of martensite formed is suppressed, and the target strength can be controlled.
In addition, the effect | action of each component is mentioned later.

上記高強度鋼板として、上記した組織を有するとともに、板厚が１．０ｍｍから３．０ｍｍ、引張強さＴＳ（ＭＰａ）が１２００ＭＰａ以上で、伸び値１３（％）以上であるものが好ましい。(ＪＩＳ５号試験片)
そのような鋼板は、上述の組織を有していて高い強度と良い伸び特性とを兼ね備えるものだからである。 The high-strength steel plate preferably has the above-described structure, a plate thickness of 1.0 mm to 3.0 mm, a tensile strength TS (MPa) of 1200 MPa or more, and an elongation value of 13 (%) or more. (JIS No. 5 specimen)
This is because such a steel sheet has the above-described structure and has both high strength and good elongation characteristics.

請求項に係る高強度鋼板の製造方法は、
1) Ｃ：０．１３〜０．２１、Ｓｉ：０．５〜２．０、Ｍｎ：０．２〜１．０
Ｃｒ：１.０〜４．０、Ｎｉ：０．０２〜１．０、 Ｍｏ：０.０５〜０．４
Ｐ：〜０．０１０、 Ｓ：〜０．００３、Ｎ：０．００５〜０．０１５
（各重量％）を含み、残部はFeおよび不可避的不純物の組成となるスラブ(圧延素材)を、
２）加熱炉抽出温度１２５０℃以上、粗圧延機出側温度を１０３０℃以上で、粗最終３パスの各々の圧下率が３０％以上とする粗圧延を行い、
３）仕上圧延機出側温度９５０℃以上、仕上前段圧延機１〜３圧延機〈６台圧延機の場合。７台圧延機の場合は１〜４圧延機〉の１台当りの圧下率４０％以上、仕上後段３圧延機の圧下の累積歪は０．５以上である仕上圧延を行い、
４）仕上圧延後２秒〜６秒の空冷をしたのちに水冷冷却し、巻き取り温度５５０℃〜６５０℃とする。 The manufacturing method of the high strength steel sheet according to the claim
1) C: 0.13-0.21, Si: 0.5-2.0, Mn: 0.2-1.0
Cr: 1.0-4.0, Ni: 0.02-1.0, Mo: 0.05-0.4
P: -0.010, S: -0.003, N: 0.005-0.015
(Each weight%), the balance is a slab (rolling material) that has a composition of Fe and inevitable impurities,
2) Rough rolling at a heating furnace extraction temperature of 1250 ° C. or higher, a rough rolling mill outlet temperature of 1030 ° C. or higher, and a rolling reduction of each of the final three passes of 30% or higher,
3) Finishing mill exit side temperature 950 ° C. or higher, pre-finishing rolling mill 1 to 3 rolling mills <6 rolling mills. In the case of 7 rolling mills, the rolling reduction is 40% or more per 1 of 1 to 4 rolling mills, and the finish rolling of the rolling mill after the finishing 3 rolling mill is 0.5 or more.
4) Air cooling is performed for 2 seconds to 6 seconds after finish rolling, and then water cooling is performed to obtain a winding temperature of 550 ° C to 650 ° C.

ホットストリップミルの熱間圧延後の急速冷却とその後巻き取り過程における温度保持の温度履歴（図１）により低合金ベイナイト均一組織を得ることにより強度の向上を目的とし、合金元素としてＣｒとＳｉを添加し、ＭｎとＮｉが少ない成分系を選択することによりマルテンサイト及び残留オーステナイトが微細分散したべイナイト均一組織を得ることができる。（図７−ｂ）
また、Ｓｉの添加により炭化物の析出を制御し、さらに均一なベイナイト組織を得ることにより炭素濃度０．８％以上のオーステナイトを多量に残留させることが可能になる。このことにより高強度で加工性が良い優れた鋼板を得ることが出来るのである。（図１１）
熱間圧延仕上げ温度を９５０℃以上の高温にすることにより、旧オーステナイト粒のアスペクト比を２．０以下にすることもできる。（図３） The purpose is to improve the strength by obtaining a low alloy bainite uniform structure by the rapid cooling after hot rolling of the hot strip mill and the temperature history of the temperature holding in the subsequent winding process (FIG. 1). By adding and selecting a component system with less Mn and Ni, a uniform structure of bainite in which martensite and retained austenite are finely dispersed can be obtained. (Fig. 7-b)
Moreover, it becomes possible to control the precipitation of carbides by addition of Si, and to obtain a more uniform austenite having a carbon concentration of 0.8% or more by obtaining a uniform bainite structure. As a result, an excellent steel sheet having high strength and good workability can be obtained. (Fig. 11)
By setting the hot rolling finishing temperature to a high temperature of 950 ° C. or higher, the aspect ratio of the prior austenite grains can be made 2.0 or less. (Figure 3)

上記の仕上圧延の際、圧延材最トップ部の圧延ロールへの噛み込み不良防止の為、圧延材の最トップ部を必要に応じ前段圧延機１〜５圧延機〈仕上圧延機６段の場合。仕上圧延機７段の場合は圧延機１〜６〉にその圧延機の予定圧下量（所定の圧延のための本来の圧下量）の１０％以下の圧下量を付加して圧下するのがよい。その圧延長さは圧延材最トップ部の噛み込みより５ｍ以内とする。
または、圧延中の圧延材と圧延ロールとのスリップ発生防止の為、仕上最終圧延機より１〜３圧延機の作業ロールに特殊ハイグリップロールを使用するとよい。
発明者らの製造試験によると、後述のように、こうした条件によって円滑に、上述の高強度鋼板を得ることができた。 In the case of the above finish rolling, in order to prevent biting failure in the rolling roll at the top of the rolled material, the top of the rolled material is used as necessary in the case of the first-stage rolling mill 1-5 rolling mill <the finishing rolling mill 6-stage. . In the case of a finish rolling mill with seven stages, it is preferable to reduce the rolling mills 1 to 6> by adding a rolling amount of 10% or less of the planned rolling amount of the rolling mill (original rolling amount for a predetermined rolling). . The rolling length is within 5 m from the biting of the top of the rolled material.
Or it is good to use a special high grip roll for the work roll of 1-3 rolling mills from a finishing final rolling mill, in order to prevent slip generation with the rolling material and rolling roll in rolling.
According to the inventors' production test, as described later, the above-described high-strength steel plate could be obtained smoothly under these conditions.

請求項に記載の高強度鋼板は、ベイナイト組織中に体積率で５％以上２０％以下の残留オーステナイトが１０μｍ平方に７個以上微細に分散した状態で混在するため、互いに相反する特性である強度と加工特性を兼備した鋼板であり、耐遅れ破壊特性にも優れた鋼板である。   The high-strength steel sheet according to the claim is a strength that is mutually contradictory because the austenite having a volume ratio of 5% or more and 20% or less is mixed in a finely dispersed state of 7 or more in 10 μm square in the bainite structure. Steel sheet with both processing characteristics and excellent delayed fracture resistance.

請求項に記載した製造方法によれば、上記した高強度鋼板を円滑に製造することが出来る。   According to the manufacturing method described in the claims, the above-described high-strength steel plate can be manufactured smoothly.

以下、１２００ＭＰａ以上の引張強度をもちながらも、優れた加工性と耐遅れ破壊特性が必要とされる加工部品に使用される薄鋼板とその製造方法について、実施の形態を示す。
鋼板の成分系として、
Ｃ：０．１３〜０．２１、Ｓｉ：０．５〜２．０、Ｍｎ：０．２〜１．０
Ｃｒ：１.０〜４．０、Ｎｉ：０．０２〜１．０、 Ｍｏ：０.０５〜０．４
Ｐ：〜０．０１０、 Ｓ：〜０．００３、Ｎ：０．００５〜０．０１５
（各重量％）を含み、残部はＦｅおよび不可避的不純物の組成とするものである。
なお、ここで述べる薄鋼板とは、板厚が１．０ｍｍから３．０ｍｍの鋼板のことである。製造する鋼板は、主として自動車、家電製品、電子機器製品等の高い加工性と強度が必要な部品に使用することが出来る。その他、鋼管用の素材として適用が可能である。 Hereinafter, embodiments of a thin steel plate used for a processed part that requires excellent workability and delayed fracture resistance while having a tensile strength of 1200 MPa or more and a manufacturing method thereof will be described.
As a component system of steel plates,
C: 0.13-0.21, Si: 0.5-2.0, Mn: 0.2-1.0
Cr: 1.0-4.0, Ni: 0.02-1.0, Mo: 0.05-0.4
P: -0.010, S: -0.003, N: 0.005-0.015
(The respective weight percentages) are included, and the balance is the composition of Fe and inevitable impurities.
In addition, the thin steel plate described here is a steel plate having a plate thickness of 1.0 mm to 3.0 mm. The steel sheet to be manufactured can be used mainly for parts that require high workability and strength, such as automobiles, home appliances, and electronic equipment products. In addition, it can be applied as a material for steel pipes.

まず、鋼板の成分について述べる。
炭素（Ｃ）としては、０．１３〜０．２１％の範囲の量が必要である。
Ｃは残留オーステナイトを安定化させるために最も重要な元素で、０．１３％未満では十分安定度が得られないので、０．１３％以上のＣ量が必要である。一方、Ｃ量が０．２１％以上になると、溶接部が硬化しすぎて溶接部から破断しやすくなる。これは、薄鋼板にとっては使用上の制約になるので、Ｃ量に上限を設けた。そして、０．１３〜０．２１％のＣ量であれば、本発明の主旨にそった複合組織が得られることを見出したものである。 First, the components of the steel sheet will be described.
As carbon (C), an amount in the range of 0.13 to 0.21% is required.
C is the most important element for stabilizing the retained austenite, and if it is less than 0.13%, sufficient stability cannot be obtained, so a C amount of 0.13% or more is necessary. On the other hand, when the C content is 0.21% or more, the welded portion is excessively hardened and is easily broken from the welded portion. Since this is a restriction in use for thin steel plates, an upper limit is set for the C amount. And it has been found that when the amount of C is 0.13 to 0.21%, a composite structure in accordance with the gist of the present invention can be obtained.

シリコン（Ｓｉ）量は、０．５〜２．０％の範囲とする。Ｓｉも残留オーステナイトの安定化のために活用する。Ｓｉは固溶強化による強度の向上効果も有する。Ｓｉ量は、０．５％以上であれば、本発明の複合組織と材質特性が得られる。Ｓｉ量は多いほど、残留オーステナイト量を増やすことができると同時に、その安定性を促す。しかし、２．０％以上のＳｉ量になると、強度延性バランス特性が飽和するので、コスト低減の観点からＳｉ量の上限を２．０％とする。   The amount of silicon (Si) is in the range of 0.5 to 2.0%. Si is also used to stabilize retained austenite. Si also has an effect of improving strength by solid solution strengthening. If the amount of Si is 0.5% or more, the composite structure and material characteristics of the present invention can be obtained. As the amount of Si increases, the amount of retained austenite can be increased and at the same time the stability thereof is promoted. However, when the Si amount is 2.0% or more, the strength ductility balance characteristic is saturated, so the upper limit of the Si amount is set to 2.0% from the viewpoint of cost reduction.

クロム（Ｃｒ）量は、１．０〜４．０％の範囲とする。Ｃｒはベイナイト組織を形成し、鋼の強度を向上させることが出来る。
Ｃｒ量が１．０％未満になると、フェライト量が多くなり高い鋼板強度が得られないので、Ｃｒ量は１．０％以上とする。Ｃｒ量が４．０％を超えると、マルテンサイトが生成しやすくなり鋼板強度が極めて高くなるとともに耐遅れ破壊性も劣化するので、その上限を４．０％とした。 The amount of chromium (Cr) is in the range of 1.0 to 4.0%. Cr forms a bainite structure and can improve the strength of the steel.
If the Cr content is less than 1.0%, the ferrite content increases and high steel sheet strength cannot be obtained, so the Cr content is 1.0% or more. If the Cr content exceeds 4.0%, martensite is easily generated, the steel sheet strength becomes extremely high and the delayed fracture resistance deteriorates, so the upper limit was made 4.0%.

マンガン（Ｍｎ）量は、０．２〜１．０％の範囲とする。Ｍｎ量が０．２%未満になると、製鋼上での製造が困難になるので０．２％以上とする。
高い強度を得るためにはＭｎを多量に添加することが好まれるが、余り高くし過ぎるとマルテンサイトが生成しやすくなり、本発明の目的とする組織が得られない。そこでＭｎ量の上限を１．０％とする。 The amount of manganese (Mn) is in the range of 0.2 to 1.0%. If the amount of Mn is less than 0.2%, it becomes difficult to manufacture on steel making, so 0.2% or more.
In order to obtain a high strength, it is preferable to add a large amount of Mn. However, if it is too high, martensite is likely to be generated, and the target structure of the present invention cannot be obtained. Therefore, the upper limit of the amount of Mn is set to 1.0%.

ニッケル（Ｎｉ）量は、０．０２〜１．０％の範囲とする。Ｎｉは固溶強化により鋼の強度を向上させることが出来るが、余り高くし過ぎるとマルテンサイトが生成しやすくなり。さらに故意に添加を行えばコストの上昇を招くため、その上限を１．０％とした。   The amount of nickel (Ni) is in the range of 0.02 to 1.0%. Ni can improve the strength of the steel by solid solution strengthening, but if it is too high, martensite is likely to be generated. Furthermore, if the addition is performed intentionally, the cost increases, so the upper limit was made 1.0%.

モリブデン（Ｍｏ）は、Ｃｒ同様にベイナイト組織を形成し、鋼の強度を向上させることが出来る。またＭｏ炭化物による水素トラップ作用を生かし耐遅れ破壊の対策とした、故意に添加を多く行えば再結晶抑制し過ぎたり、コストの上昇を招くため、その範囲を０．０５〜０．４０％とした。   Molybdenum (Mo) can form a bainite structure like Cr, and can improve the strength of steel. In addition, taking advantage of the hydrogen trap action by Mo carbides as a countermeasure for delayed fracture resistance, intentionally adding too much will suppress recrystallization too much or cause an increase in cost, so the range is 0.05 to 0.40%. did.

燐（Ｐ）は、溶接性向上のため出来るだけ少なくすることが必要で、その上限の量を０．０１０％とした。   Phosphorus (P) needs to be reduced as much as possible in order to improve weldability, and the upper limit amount is set to 0.010%.

硫黄（Ｓ）も、溶接性向上のため出来るだけ少なくすることが必要で、その上限の量を０．００３％とした。   Sulfur (S) is also required to be reduced as much as possible in order to improve weldability, and the upper limit amount is set to 0.003%.

窒素（Ｎ）量は、０．００５〜０．０１５％の範囲とする。窒素は、炭素と同様にオーステナイト相安定化元素であるが、多すぎると溶接性を低下させるため、その範囲を０．００５〜０．０１５％とした。   The amount of nitrogen (N) is in the range of 0.005 to 0.015%. Nitrogen is an austenite phase stabilizing element like carbon, but if it is too much, the weldability is lowered, so the range was made 0.005 to 0.015%.

上記の基準成分に調整したスラブは、再加熱してから熱間圧延をおこなうか、もしくは鋳造後直ちに熱間圧延をおこなうものとする。
図１は、この発明の実施形態の製造プロセスにおける熱間圧延での温度履歴の概念と旧オーステナイト粒径を示すもので、横軸は時間経過、縦軸は温度である。
熱間圧延を施すにあたっては、加熱炉抽出温度を１２５０℃以上とした。これは仕上後面温度９５０℃確保を最優先とし、その為に加熱炉でオーステナイト粒が大きくなってもやむをえないとした。但し圧延工程でオーステナイト粒を小さくする。その為に、仕上圧延機前に、極力旧オーステナイトを細粒化する必要がある。そこで粗圧延において、粗圧延機出側温度を１０３０℃以上で、粗最終３パスの各々の圧下率が３０％以上にすることにより、結晶粒を３５μｍ以下とする。図２は仕上げ前クロップシャーでカットした粗圧延後の旧オーステナイト粒径を示す。 The slab adjusted to the above-described reference component is either hot-rolled after reheating or hot-rolled immediately after casting.
FIG. 1 shows the concept of temperature history in hot rolling and the prior austenite grain size in the manufacturing process according to the embodiment of the present invention. The horizontal axis represents time and the vertical axis represents temperature.
In performing hot rolling, the furnace extraction temperature was set to 1250 ° C. or higher. For this reason, securing the finishing surface temperature of 950 ° C. was given top priority, and for that reason, it was inevitable that the austenite grains became large in the heating furnace. However, austenite grains are reduced in the rolling process. Therefore, it is necessary to refine the old austenite as much as possible before the finishing mill. Therefore, in rough rolling, the temperature of the roughing mill is 1030 ° C. or higher, and the rolling reduction of each of the final three passes is 30% or higher, so that the crystal grains are 35 μm or lower. FIG. 2 shows the prior austenite grain size after rough rolling cut by a pre-finishing crop shear.

仕上前段の第１〜３段圧延機〈仕上６段圧延機の場合。仕上７台圧延機の場合は第１〜４段圧延機〉の１台当りの圧下率４０％以上。 仕上後段３圧延機の圧下の累積歪は０．５以上で、仕上圧延機出側温度として９５０℃以上を確保して、オーステナイト粒径を１０μｍ以下とする。仕上圧延後２〜６秒は空冷をし、その後水冷冷却とする、巻き取り温度５５０℃〜６５０℃とするが、前述の空冷でオーステナイト粒の整粒化をはかる。即ち熱間圧延にあたっては、熱間圧延後ホットラン冷却が開始する前までに旧オーステナイトの大きさを１０μｍ以下にし、且つその旧オーステナイト粒は加工歪のない整粒化されたものとした。
図３は、ＳＥＭ組織観察による本発明鋼の旧オーステナイト粒の観察結果である。旧オーステナイト粒の平均粒径は９．３μｍで均一な整粒組織を呈していて、その長軸/短軸の平均アスペクト比は１．７である。 1st to 3rd rolling mills before finishing (in the case of a finishing 6-high rolling mill. In the case of a finishing 7-roll mill, the rolling reduction per unit is 40% or more. The cumulative strain under rolling of the finishing third rolling mill is 0.5 or more, and the finishing mill exit temperature is ensured to be 950 ° C. or more and the austenite grain size is set to 10 μm or less. Air cooling is performed for 2 to 6 seconds after the finish rolling, and then the cooling temperature is set to 550 ° C. to 650 ° C. The austenite grains are sized by air cooling as described above. That is, in the hot rolling, the size of the prior austenite was set to 10 μm or less before the hot run cooling was started after the hot rolling, and the prior austenite grains were sized without processing strain.
FIG. 3 is a result of observation of prior austenite grains of the steel of the present invention by SEM structure observation. The prior austenite grains have an average particle size of 9.3 μm and a uniform sized structure, and the average aspect ratio of the major axis / minor axis is 1.7.

圧延完了温度が９５０℃以下の低温度圧延で、仕上後段３圧延機の累積歪が０．５以下の場合にはオーステナイト粒が大きくなり（１０μｍ以上）、しかもオーステナイト粒の形が圧延された扁平となり、異方性の大きな原因となる。 図４は仕上圧延機出側温度（ＦＤＴ）と伸びの異方性の関係をしめす。ＦＤＴが９５０℃以下になると伸びの異方性が現れてくる。なお、異方性は｜Ｃ−Ｌ｜/(Ｃ+Ｌ)/２で定義した（Ｌは圧延方向の伸び、Ｃはそれと直角な方向の伸びである）。値が小さいほど異方性が少ないことを示している。
ここで「歪み」とは、各スタンド（各段、または粗圧延時の各パス）の入側での鋼板の厚さｈ0と出側での厚さｈ1の差を両者の平均厚さで除した
ε=（ｈ0−ｈ1）／｛（ｈ0＋ｈ1）／２｝
をいい、「累積歪み」とは、後段３スタンドの各段（各パス）での歪みを金属組織に対する影響の強さを考慮して加重積算したもので、最終段（最終パス）とその前段（前パス）・前々段（前々パス）での歪みをそれぞれεn、εn-1、εn-2とするとき、
εC＝εn＋εn-1／２＋εn-2／４
で表されるεCをいうものとする。 In the case of low temperature rolling with a rolling completion temperature of 950 ° C. or lower and the cumulative strain of the final stage 3 rolling mill is 0.5 or less, the austenite grains become large (10 μm or more) and the shape of the austenite grains is rolled. This is a major cause of anisotropy. FIG. 4 shows the relationship between finishing mill exit temperature (FDT) and elongation anisotropy. When FDT is 950 ° C. or lower, anisotropy of elongation appears. The anisotropy was defined by | CL | / (C + L) / 2 (L is the elongation in the rolling direction, and C is the elongation in the direction perpendicular thereto). A smaller value indicates less anisotropy.
Here, “strain” means the difference between the thickness h 0 of the steel sheet on the entry side of each stand (each stage or each pass during rough rolling) and the thickness h 1 on the delivery side, and the average thickness of the two. Ε = (h 0 −h 1 ) / {(h 0 + h 1 ) / 2} divided by
"Cumulative strain" is a weighted integration of the strain at each stage (each pass) of the latter three stands in consideration of the strength of the influence on the metal structure. The last stage (final path) and its preceding stage (Previous pass) and when the distortion in the previous stage (previous pass) is ε n , ε n-1 and ε n-2 , respectively.
ε C = ε n + ε n-1 / 2 + ε n-2 / 4
Ε C represented by

高温仕上げ圧延を行う為に、圧延での加工発熱による鋼板の温度上昇を利用する、そのために、後段圧延機の高歪圧延スケジュールはもちろんのこと、前段スタンドの圧下率も４０％以上とすることが大事である。図５に示した様に、粗厚に相違があるものの圧下率により同じ圧延サイズで、鋼種によっては仕上後面温度が８０℃違うことがわかる。   In order to perform high-temperature finish rolling, the temperature rise of the steel sheet due to processing heat generated during rolling is used. Therefore, not only the high strain rolling schedule of the latter rolling mill but also the rolling reduction of the former stand should be 40% or more. Is important. As shown in FIG. 5, it can be seen that although the thickness is different, the finishing surface temperature is different by 80 ° C. depending on the steel type with the same rolling size depending on the rolling reduction.

９５０℃以上で熱間圧延を完了し、仕上圧延後ホットラン冷却を行なわず空冷の時間を２〜６秒取り、結晶粒中の転位密度を減少させる。図６に、同一鋼種について圧延温度を変えたときの計算による仕上Ｆ１圧延機からホットラン冷却開始までのオーステナイト粒径の変化と転位密度の変化を示す。この図より、転位密度は圧延温度に大きく影響を受けることがわかる。また、オーステナイト粒径は高圧下圧延条件下では、Ａｒ3変態以上の温度であれば低い温度の方が小さくなることもわかる。ただし、転位密度は高くなり、異方性の高い材料となる。圧延後のホットランの空冷で転位密度が大きく減少しているが、６秒以内が効果的であることがわかる。ここで、アスペクト比を２．０以下にするということは、このシミュレーションモデルでは、転位密度を少なくとも２．５０Ｅ＋１０（ρ／ｃｍ2）以下にすることとなる（実績とモデルとの付き合せ結果）。しかしながら転位密度の減少は旧オーステナイト粒径を大きくすることとなる。ここで旧オーステナイト粒径を１０μｍ以下で転位密度を上記の数値以下にするためには、上記圧延条件（圧延温度：９５０℃以上、空冷時間：２〜６秒）が必要となる。
なお、今回適用のシミュレーションモデルは、柳本、森本らの「鉄と鋼」Vol.88(2002）No.11をベースとして、数式の係数は今回用に再構築したものである。 Hot rolling is completed at 950 ° C. or higher, and after the finish rolling, hot run cooling is not performed and air cooling is performed for 2 to 6 seconds to reduce the dislocation density in the crystal grains. FIG. 6 shows changes in the austenite grain size and changes in the dislocation density from the finish F1 rolling mill to the start of hot-run cooling by calculation when the rolling temperature is changed for the same steel type. From this figure, it can be seen that the dislocation density is greatly influenced by the rolling temperature. It can also be seen that the austenite grain size becomes smaller at a temperature lower than the Ar3 transformation under high-pressure rolling conditions. However, the dislocation density increases and the material becomes highly anisotropic. Although the dislocation density is greatly reduced by air cooling of the hot run after rolling, it can be seen that the effect is within 6 seconds. Here, setting the aspect ratio to 2.0 or less means that in this simulation model, the dislocation density is set to at least 2.50E + 10 (ρ / cm 2 ) or less (result of matching the results with the model). . However, the reduction in dislocation density increases the prior austenite grain size. Here, the above rolling conditions (rolling temperature: 950 ° C. or higher, air cooling time: 2 to 6 seconds) are required in order to make the prior austenite grain size 10 μm or less and the dislocation density below the above numerical value.
The simulation model applied this time is based on Yanagimoto, Morimoto et al. "Iron and Steel" Vol.88 (2002) No.11, and the coefficients of the formula are reconstructed for this time.

巻き取り温度を５５０℃以上、６５０℃以下としたが、５５０℃以下の温度範囲ではマルテンサイト組織が多く生成し、遅れ破壊がおこりやすい。また６５０℃以上の温度ではフェライトとパーライト組織が生成し、高い強度が得られない。図７に３種類の高強度鋼板の断面組織を示す。
いずれもベイナイト基地で、a)はマルテンサイトが多い組織、ｂ）はマルテンサイトが少なく微細な組織、c)はフェライトを含む組織の写真を示す。ｂ）が本発明の組織である。 The coiling temperature is set to 550 ° C. or more and 650 ° C. or less. However, in the temperature range of 550 ° C. or less, a lot of martensite structure is generated and delayed fracture is likely to occur. Further, at a temperature of 650 ° C. or higher, ferrite and pearlite structure are formed, and high strength cannot be obtained. FIG. 7 shows cross-sectional structures of three types of high-strength steel plates.
Both are bainite bases, a) shows a structure with a lot of martensite, b) shows a fine structure with little martensite, and c) shows a picture of a structure containing ferrite. b) is the organization of the present invention.

ベイナイト組織では、旧オーステナイト粒界を始めとしてパケット境界やブロック境界にも、すなわち旧オーステナイト粒内にも、オーステナイトが残留する。この残留オーステナイトは、ベイナイト組織を母相としかつ変態前の旧オーステナイト粒径を１０μｍ以下にすることで、粒径１μｍ以下と極めて微細で、１０μｍ平方に７個以上と緻密で均一に分散させることが出来る。図８は、ＥＢＳＰ法を用いて、体心立方構造のベイナイト相と面心立方構造のオーステナイト相を色分けした本発明鋼の組織断面を示した。明るい色で示した残留オーステナイト組織は１．０μｍ以下で１０μｍ平方に７個以上の微細かつ均一に分散している。
これらの熱間圧延の制御で残留オーステナイトが微細かつ均一に分散したベイナイト組織を得るのである。 In the bainite structure, austenite remains not only in the former austenite grain boundary but also in the packet boundary and the block boundary, that is, in the old austenite grain. This retained austenite has a bainite structure as the parent phase and the prior austenite grain size before transformation is 10 μm or less, so that it is extremely fine with a grain size of 1 μm or less, and is finely and evenly dispersed with 7 or more in 10 μm square. I can do it. FIG. 8 shows a structural cross section of the steel of the present invention in which the bainite phase having a body-centered cubic structure and the austenite phase having a face-centered cubic structure are color-coded using the EBSP method. The retained austenite structure shown in a bright color is finely and uniformly dispersed in 1.0 μm or less and 7 μm or more in 10 μm square.
By controlling the hot rolling, a bainite structure in which retained austenite is finely and uniformly dispersed is obtained.

なお、高強度で板厚が薄い材料〈板厚２ｍｍ以下〉を高歪・高圧化率圧延すると、圧延の際に、板トップ部のかみ込み性不良や圧延中にロールと圧延材との間でスリップが発性し易くなる。これまでの圧延結果では、圧延材の最トップ部の噛み込み性はＴＳ<１０００ＭＰａクラスでは各圧延機一台当りの圧下率４０〜５０％の圧延では問題ないがTＳ>１０００ＭＰａクラスの材料となると最終圧延機、及び前圧延機の１〜２圧延機で圧延材トップ部の噛み込み不良が多発（発生率５０％）し出す。この対策としてロール摩擦係数を上げる為、ロール研削仕上がり粗度Ｒａを１μｍ〈通常０．５μｍ〉まで上げて、圧延中の摩擦係数μを０．４〈通常０．３〉に上げる、又圧延材最トップ部の温度を下げない為にロール冷却水を絞る等の対策を施したが、確実な効果か得られなかった。そこで、図９に示す如く圧延機出側の圧延材最トップ部を長さ５ｍ以内で、薄め板厚（予定板厚の１０位薄めの板厚）とし、それより予定板厚迄の傾斜を付けることとした。
これに依り、噛み込み不良が激減した（発生率０）。仕上前段圧延機より仕上最終圧延機前圧延機迄の圧下設定量を予定設定量の１０％以内を加算した設定とする。そしてその圧下設定時間は板トップ部の圧延機噛み込みから２秒以内とする。 In addition, if a high-strength material with a thin plate thickness (plate thickness of 2 mm or less) is rolled at a high strain / high pressure ratio, during the rolling, the bite at the top of the plate will be poor, or the roll and the rolled material will not be rolled during rolling. This makes slipping easily. According to the rolling results so far, the biting property of the top part of the rolled material is TS <1000 MPa class, there is no problem in rolling with a rolling reduction of 40 to 50% per rolling mill, but it becomes a material of TS> 1000 MPa class. In the final rolling mill and the 1-2 rolling mills of the previous rolling mill, biting defects at the top of the rolled material frequently occur (occurrence rate 50%). In order to increase the roll friction coefficient as a countermeasure, the roll grinding finish roughness Ra is increased to 1 μm (usually 0.5 μm), and the friction coefficient μ during rolling is increased to 0.4 (normally 0.3). Although measures such as squeezing roll cooling water were taken to prevent the temperature of the topmost part from being lowered, a reliable effect could not be obtained. Therefore, as shown in FIG. 9, the top part of the rolled material on the delivery side of the rolling mill is within 5 m in length, and is made a thin plate thickness (a plate thickness that is tenth thinner than the planned plate thickness), and then the inclination to the planned plate thickness is increased. I decided to attach it.
As a result, biting defects were drastically reduced (occurrence rate 0). The reduction setting amount from the finishing pre-rolling mill to the finishing final rolling mill pre-rolling mill is set to a value obtained by adding up to 10% of the preset setting amount. The reduction setting time is set to be within 2 seconds from the biting of the rolling mill at the plate top portion.

圧延中のロールと圧延材のスリップについて
最終板厚２ｍｍ以下でＴＳ>１０００ＭＰａクラスの材料を高温、高圧下率圧延すると、最終圧延機とその前圧延機でスリップが発生し易くなる。「現象としては圧延中に突然金属音が発し、スリップ発生時の圧延機の圧延荷重が急激に５０％近くにも落ちる、圧延ロールが空転し、圧延板が前進しなくなる。この時の圧延ロールを圧延機より引き抜きロール粗度を測定すると、ロール粗度Ｒａで０．１μｍ以下となり圧延材とロールがスリップしやすい状況となった。そこで、この対策として、特殊ハイグリップロールを使用した。この結果スリップ事故は皆無となった。このロールは微小炭化物（１μｍ以下）をロール表面全体に均一分散させて、その炭化物を一種のスパイクにしてそれを硬い基地でささえる。また微小炭化物がなくなっても、その下から次の微小酸化物が出てきて安定した高摩擦係数が維持出来、スリップ発生はなくなる。図１０に示した通り、圧延による摩擦係数の変化は一般ロールと比較して安定した摩擦係数（０．３）を維持している。 About slip of roll and rolled material during rolling When a material of TS> 1000 MPa class is rolled at a high temperature and high pressure under a final sheet thickness of 2 mm or less, slip is likely to occur in the final rolling mill and the preceding rolling mill. “As a phenomenon, a metal noise is suddenly generated during rolling, and the rolling load of the rolling mill at the time of slipping suddenly drops to nearly 50%, the rolling roll is idled, and the rolling plate does not move forward. When the roll roughness of the drawn roll was measured from a rolling mill, the roll roughness Ra was 0.1 μm or less, and the rolled material and the roll were likely to slip, so a special high grip roll was used as a countermeasure. As a result, there was no slip accident.This roll uniformly disperses fine carbide (less than 1 μm) over the roll surface, and makes the carbide a kind of spike and holds it on a hard base. Then, the next minute oxide comes out from below, a stable high friction coefficient can be maintained, and no slip is generated.As shown in FIG. This change maintains a stable friction coefficient (0.3) as compared with a general roll.

図１１は、図１の製造プロセスにより製造した熱延鋼板の残留オーステナイトの体積率Vγと引張試験との関係を示したものである。ａ)は体積率Ｖγと引張強さ×伸びの関係を示し。ｂ)は体積率Ｖγと伸びの関係を示す。残留オーステナイトの体積率Ｖγが５から２０％の範囲で体積率Ｖγが多くなるにしたがって引張強さ×伸びおよび伸びが改善されていることが示されている。なお、このデータの金属組織は図７−ｂのマルテンサイトが少なく微細なベイナイト組織をしめしている。
本発明は、以上の知見に基づき開発されたものである。 FIG. 11 shows the relationship between the volume ratio Vγ of retained austenite of the hot-rolled steel sheet manufactured by the manufacturing process of FIG. 1 and the tensile test. a) shows the relationship between the volume fraction Vγ and the tensile strength × elongation. b) shows the relationship between the volume fraction Vγ and elongation. It is shown that the tensile strength × elongation and elongation are improved as the volume fraction Vγ increases in the range of the volume fraction Vγ of retained austenite from 5 to 20%. The metal structure of this data shows a fine bainite structure with few martensites as shown in FIG.
The present invention has been developed based on the above findings.

以下に発明の実施例を説明する。
表１に示す化学成分を有する溶鋼を、連続鋳造法もしくは鍛造法によりスラブ（圧延素材）とした。続いてこれらのスラブを再加熱し、熱間圧延を行い、熱延鋼板とした。表２に熱間圧延条件とその材料特性を示す。 Examples of the invention will be described below.
The molten steel having chemical components shown in Table 1 was used as a slab (rolled material) by a continuous casting method or a forging method. Subsequently, these slabs were reheated and hot-rolled to obtain hot-rolled steel sheets. Table 2 shows the hot rolling conditions and the material properties.

表１に示す鋼種Ａ、Ｂ、Ｃは本発明の範囲で、Ｄ、Ｅ、Ｆ、Ｇ、Ｈ、Ｉは比較例である。   Steel types A, B, and C shown in Table 1 are within the scope of the present invention, and D, E, F, G, H, and I are comparative examples.

比較例の鋼種ＤはＳｉが低く、かつＮｉが高く本発明の範囲から外れたものである。
比較例の鋼種ＥはＳｉが低く本発明の範囲から外れたものである。
鋼種Ｆ、ＧはＣが低く本発明範囲から外れたもので、鋼種ＩはＣが高く本発明範囲から外れたものである。鋼種ＨはＣｒが高く本発明範囲から外れたものである。 Steel type D of the comparative example is low in Si, high in Ni, and out of the scope of the present invention.
Steel type E of the comparative example has low Si and is out of the scope of the present invention.
Steel types F and G have low C and deviated from the scope of the present invention, and steel type I has high C and deviated from the scope of the present invention. Steel type H is high in Cr and is out of the scope of the present invention.

表２のＮｏ.１〜６は成分範囲を満足する鋼種Ａ、ＢおよびＣを用いて圧延を行った事例である。
Ｎｏ.１はＳｉが０．５１％である鋼種Aを用いて熱間圧延巻き取り温度を６５５℃としたものである。この場合ＴＳ×ＥＬはきわめて良好であるが引張強さが７７８ＭＰａときわめて低い。
Ｎｏ.２は鋼種Ａを用いて熱間圧延巻き取り温度を６３０℃としたものである。引張強さは１２００ＭＰａ以上あり、伸びも１３％以上あり良好である。
Ｎｏ.３はＳｉが１．００％である鋼種Ｂを用いて熱間圧延巻き取り温度を５９５℃としたものである。強度および伸びともに良好であり、Ｎｏ.２よりもなお強度および伸びともに向上している。
Ｎｏ．４およびＮｏ．５は熱間圧延の圧下率を満足しておらず、強度・伸びは満足しているが遅れ破壊が発生している。
Ｎｏ.６はＳｉが１．４４％である鋼種Ｃを用いて熱間圧延巻き取り温度を６１０℃としたものである。強度および伸びともに良好であり、Ｎｏ．３よりもなお強度および伸びともに向上している。 Nos. 1 to 6 in Table 2 are examples of rolling using steel types A, B, and C that satisfy the component ranges.
No. 1 uses a steel type A in which Si is 0.51% and the hot rolling coiling temperature is 655 ° C. In this case, TS × EL is very good, but the tensile strength is very low at 778 MPa.
No. 2 uses steel type A and the hot rolling coiling temperature is 630 ° C. The tensile strength is 1200 MPa or more, and the elongation is 13% or more.
No. 3 uses steel type B with Si of 1.00% and the hot rolling coiling temperature is 595 ° C. Both strength and elongation are good, and both strength and elongation are higher than those of No.2.
No. 4 and no. No. 5 does not satisfy the rolling reduction ratio of hot rolling, and the strength and elongation are satisfied, but delayed fracture occurs.
No. 6 is a steel type C in which Si is 1.44% and the hot rolling coiling temperature is 610 ° C. Both strength and elongation are good. Both strength and elongation are improved more than 3.

Ｎｏ．７〜１２は本発明の成分範囲を外れた比較例の鋼種を用いて、熱間圧延をおこなったものである。   No. Nos. 7 to 12 are hot-rolled using the steel types of comparative examples that are out of the component range of the present invention.

Ｎｏ．７はＳｉが低く、Ｎｉが高い鋼種Ｄを用いて圧延したものである。スポット溶接性（Ｓ／Ｗ性）および遅れ破壊特性が悪い。
Ｎｏ．８はＳｉが低い鋼種Ｅを用いたもので、強度が低く、強度・延性バランスも悪い。
Ｎｏ．９，Ｎｏ．１０はＣが低い鋼種Ｆ，Ｇを用いたもので、強度が低く、強度・延性バランスも悪い。
Ｎｏ．１１，Ｎｏ．１２はＣが高い鋼種Ｈ，Ｉを用いたもので、強度は高く、強度・延性バランスもよいが、スポット溶接性および遅れ破壊特性が悪い。 No. No. 7 is rolled using a steel type D having low Si and high Ni. Spot weldability (S / W property) and delayed fracture characteristics are poor.
No. No. 8 is a steel type E with low Si, which has a low strength and a poor balance between strength and ductility.
No. 9, no. No. 10 uses steel types F and G having a low C, and has a low strength and a poor balance between strength and ductility.
No. 11, no. No. 12 uses steel types H and I with high C, which has high strength and good balance between strength and ductility, but has poor spot weldability and delayed fracture characteristics.

フェライト粒の体積率は、鋼板の圧延方向断面を研磨後、ナイタル腐食後、光学顕微鏡により観察し、市販の画像解析装置も用いて測定した。
マルテンサイトの体積率は、鋼板の圧延方向断面を研磨後、４％ピクリン酸アルコールと２％ピロ硫酸ナトリウムを１対１に混合した液でエッチングし、板厚方向１／４の位置を光学顕微鏡により観察し、画像解析処理により白色にエッチングされたマルテンサイトを測定して求めた。
残留オーステナイトの測定はＣｕのＫα線を用いてＸ線回折法により求めた。板厚１／２ｔ部位で表面電解研磨仕上げ後、オーステナイト相の（２００），（２２０）および（３１１）面とフェライト相の（２００），（２１１）面の積分強度を測定し、それぞれの組合わせから算出される残留オーステナイト体積率の平均値を用いた。
引張り特性（引張り強さＴＳ、伸び値ＥＬ）はＪＩＳ５号試験片形状にて引張試験し測定した。
遅れ破壊性は、８％余歪を負荷した引張試験片稼動部の中央部に、φ１０ｍｍのパンチ穴をクリアランス１２．２％であけた後、１規定の塩酸に浸漬後経過観察し確認した。 The volume fraction of the ferrite grains was measured by using a commercially available image analysis apparatus after polishing a cross section in the rolling direction of the steel sheet, observing with a light microscope after nitral corrosion.
The volume ratio of martensite is determined by polishing a cross section in the rolling direction of the steel sheet, etching with a mixture of 4% picric alcohol and 2% sodium pyrosulfate in a 1: 1 ratio, and measuring the position in the thickness direction 1/4 with an optical microscope. And martensite etched in white by image analysis processing was measured and determined.
The residual austenite was measured by an X-ray diffraction method using Cu Kα rays. After the surface electrolytic polishing finish at the 1/2 t thickness, the integrated strengths of the (200), (220) and (311) planes of the austenite phase and the (200) and (211) planes of the ferrite phase were measured. The average value of the retained austenite volume fraction calculated from the combination was used.
Tensile properties (tensile strength TS, elongation value EL) were measured by a tensile test using a JIS No. 5 test piece shape.
The delayed fracture property was confirmed by observing the follow-up after immersing in 1 N hydrochloric acid after punching a punch hole of φ10 mm with a clearance of 12.2% in the center of the tensile test piece working portion loaded with 8% surplus strain.

以上のように低合金組成において高強度で高延性な特性を示す実施例の高強度鋼板は、自動車構造用部材等として使用するのに好適である。
例えば自動車のセンターピラーのように、ドアの支持とともに衝突時の変形防止等に必要な引張り強度が求められる他、プレス成形等のため曲げ、絞り加工性、関連機器の取付け穴を形成するための穴拡げ加工性、さらには他の車体部品と接合するための溶接性などに高いレベルが要求される部材として、極めて好ましい鋼板といえる。 As described above, the high-strength steel sheets of the examples showing high strength and high ductility characteristics in a low alloy composition are suitable for use as automobile structural members and the like.
For example, the center pillar of an automobile requires the tensile strength necessary to prevent deformation at the time of collision as well as the support of the door, as well as bending, drawing workability for press molding, etc., to form mounting holes for related equipment It can be said that it is a very preferable steel plate as a member that requires a high level of hole expandability and further weldability for joining with other body parts.

この発明の実施形態の製造プロセスにおける熱間圧延での温度履歴の概念を示す線図。The diagram which shows the concept of the temperature history in the hot rolling in the manufacturing process of embodiment of this invention. 仕上げ入り口クロップの旧オーステナイト粒写真。Old austenite grain photograph of finishing entrance crop. 旧オーステナイト粒の写真。A photo of old austenite grains. 圧延仕上げ温度と伸びの異方性との関係。Relationship between rolling finishing temperature and elongation anisotropy. 圧延スケジュールと温度。Rolling schedule and temperature. 転位密度と旧オーステナイト粒径。Dislocation density and prior austenite grain size. 代表的な断面組織写真。Representative cross-sectional structure photograph. 成分及び圧延条件が本発明の範囲において製造した鋼板のＥＢＳＰ法による断面組織写真で、明るい（薄い）色が残留オーステナイト。A bright (light) color is a retained austenite in the cross-sectional structure photograph by the EBSP method of the steel plate which the component and rolling conditions manufactured in the range of this invention. 最トップ咬みこみ改善。Improved top bite. ロールの種類による圧下率と摩擦係数。Rolling ratio and friction coefficient depending on the type of roll. 強度・延性バランス及び残留オーステナイト量の関係。Relationship between strength and ductility balance and retained austenite amount.

Claims (7)

残留オーステナイト粒が１μｍ以下の大きさであり１０μｍ平方に７個以上分散されており、残留オーステナイト組織の体積率が５％以上２０％以下で、マルテンサイト組織の体積率が０％以上１０％以下で、残部がベイナイト組織であることを特徴とする高強度熱延鋼板。   Residual austenite grains are 1 μm or less in size and 7 or more are dispersed in a 10 μm square, the volume ratio of the retained austenite structure is 5% or more and 20% or less, and the volume ratio of the martensite structure is 0% or more and 10% or less. And the remainder is a bainite structure, The high-strength hot-rolled steel sheet characterized by the above-mentioned. 旧オーステナイト粒径が１０μｍ以下で、その平均アスペクト比が２．０以下であることを特徴とする請求項１に記載の高強度熱延鋼板。   The high-strength hot-rolled steel sheet according to claim 1, wherein the prior austenite grain size is 10 µm or less and the average aspect ratio is 2.0 or less. Ｃ：０．１３〜０．２１、Ｓｉ：０．５〜２．０、Ｍｎ：０．２〜１．０、Ｃｒ：１.０〜４．０、Ｎｉ：０．０２〜１．０、 Ｍｏ：０.０５〜０．４、Ｐ：〜０．０１０、 Ｓ：〜０．００３、Ｎ：０．００５〜０．０１５を含み、残部はＦｅおよび不可避的不純物の組成であることを特徴とする請求項１または２に記載の高強度熱延鋼板。   C: 0.13-0.21, Si: 0.5-2.0, Mn: 0.2-1.0, Cr: 1.0-4.0, Ni: 0.02-1.0, Mo: 0.05-0.4, P: -0.010, S: -0.003, N: 0.005-0.015, the balance being the composition of Fe and inevitable impurities The high-strength hot-rolled steel sheet according to claim 1 or 2. 板厚が１．０〜３．０ｍｍで、引張強さが１２００ＭＰａ以上である請求項１〜３のいずれかに記載の高強度熱延鋼板。   The high-strength hot-rolled steel sheet according to any one of claims 1 to 3, having a plate thickness of 1.0 to 3.0 mm and a tensile strength of 1200 MPa or more. Ｃ：０．１３〜０．２１、Ｓｉ：０．５〜２．０、Ｍｎ：０．２〜１．０、Ｃｒ：１.０〜４．０、Ｎｉ：０．０２〜１．０、 Ｍｏ：０.０５〜０．４、Ｐ：〜０．０１０、 Ｓ：〜０．００３、Ｎ：０．００５〜０．０１５を含み、残部はＦｅおよび不可避的不純物の組成となる鋼材を、
加熱炉抽出温度が１２５０℃以上、粗圧延機出側温度が１０３０℃以上、粗最終３パスの各々の圧下率が３０％以上となる条件で粗圧延し、
仕上圧延機の出側温度が９５０℃以上、仕上前段での１台当りの圧下率が４０％以上、仕上後段３圧延機での圧下の累積歪が０．５以上となる条件で仕上圧延し、
仕上圧延後は２秒〜６秒の空冷をしたのちに水冷冷却し、巻き取り温度を５５０〜６５０℃とする
ことを特徴とする高強度熱延鋼板の製造方法。 C: 0.13-0.21, Si: 0.5-2.0, Mn: 0.2-1.0, Cr: 1.0-4.0, Ni: 0.02-1.0, Mo: 0.05-0.4, P: -0.010, S: -0.003, N: 0.005-0.015, and the balance is a steel material having a composition of Fe and inevitable impurities,
Rough rolling under conditions where the heating furnace extraction temperature is 1250 ° C. or higher, the roughing mill outlet temperature is 1030 ° C. or higher, and the rolling reduction of each of the final three passes is 30% or higher,
Finish rolling under conditions where the exit temperature of the finishing mill is 950 ° C or higher, the rolling reduction per unit in the pre-finishing stage is 40% or more, and the cumulative strain of rolling in the post-finishing three rolling mill is 0.5 or more. ,
A method for producing a high-strength hot-rolled steel sheet, characterized in that after finish rolling, air cooling is performed for 2 seconds to 6 seconds, followed by water cooling and a winding temperature of 550 to 650 ° C. 仕上圧延の際、圧延材最トップ部の圧下量を、最終段以外の圧延機のいずれか１以上において予定圧下量より多くし、その圧下量はその圧延機の予定圧下量の１０％以下にし、その圧延長さは圧延材最トップ部の噛み込みより５ｍ以内としてその後に予定圧下量に戻すことを特徴とする請求項５に記載の高強度熱延鋼板の製造方法。   At the time of finish rolling, the reduction amount of the topmost portion of the rolled material is made larger than the planned reduction amount in any one or more of the rolling mills other than the final stage, and the reduction amount is set to 10% or less of the planned reduction amount of the rolling mill. The method for producing a high-strength hot-rolled steel sheet according to claim 5, wherein the rolling length is set to be within 5 m from the biting of the topmost portion of the rolled material, and is then returned to the planned reduction amount. 最終圧延機を含む仕上後段圧延機の作業ロールに、微小炭化物を表面に分散したハイグリップロールを使用することを特徴とする請求項５または６に記載の高強度熱延鋼板の製造方法。   The method for producing a high-strength hot-rolled steel sheet according to claim 5 or 6, wherein a high-grip roll having fine carbide dispersed on the surface is used as a work roll of a finishing post-rolling mill including a final rolling mill.
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