JP2014177669A - Non-heat treated low yield ratio high tensile thick steel plate, and manufacturing method therefor - Google Patents

Non-heat treated low yield ratio high tensile thick steel plate, and manufacturing method therefor Download PDF

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JP2014177669A
JP2014177669A JP2013051407A JP2013051407A JP2014177669A JP 2014177669 A JP2014177669 A JP 2014177669A JP 2013051407 A JP2013051407 A JP 2013051407A JP 2013051407 A JP2013051407 A JP 2013051407A JP 2014177669 A JP2014177669 A JP 2014177669A
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Takao Akatsuka
隆男 赤塚
Akio Omori
章夫 大森
Nobuyuki Ishikawa
信行 石川
Yoshi Nakagawa
佳 中川
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JFE Steel Corp
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Abstract

PROBLEM TO BE SOLVED: To provide a low yield ratio high tensile thick steel plate for building construction member excellent in earthquake resistance and a manufacturing method therefor.SOLUTION: There is provided a thick steel plate containing a component composition consisting of, by mass%, C:0.05 to 0.16%, Si:0.05 to 0.45%, Mn:1.2 to 1.8%, P, S, Al:0.05% or less, N:0.0040% or less, Ti:0.005 to 0.020%, 4.0≥Ti/N≥2.0, Ceq:0.35 to 0.48, and one or more kind of Cu, Ni, Cr, V, Ca and REM as needed and the balance Fe with inevitable impurities, and a hard phase consisting of ferrite having an average crystal particle diameter of 4.0 to 18.0 μm, one or more kinds of pearlite, bainite and martensite, and having a micro structure having a ferrite area ratio of 45% to 70%, and an average hardness of a steel plate surface part of 225 HV or less and difference of hardness between the surface part and a plate thickness center part of 60 HV or less. The steel having the component composition is applied to a two-stage cooling after hot rolling.

Description

本発明は、耐震性を必要とする建築構造部材用として好適な、非調質低降伏比高張力厚鋼板およびその製造方法に係り、特に、板厚:19mm以上で、円形鋼管柱あるいは角形鋼管柱など、冷間曲げ加工を施される用途に好適なものに関する。   The present invention relates to a non-tempered low yield ratio high-tensile steel plate suitable for building structural members that require earthquake resistance, and a method for manufacturing the same, and in particular, with a plate thickness of 19 mm or more, a circular steel pipe column or a square steel pipe. The present invention relates to a column and the like that are suitable for applications that are subjected to cold bending.

近年、地震時の安全性確保の観点から建築構造物などにおいては、素材として、優れた耐震性を有する鋼板(鋼材)を用いることが要求されている。また、これまでの研究成果で、降伏比の低い鋼板(鋼材)ほど耐震性に優れることが明らかにされている。   In recent years, from the viewpoint of ensuring safety during an earthquake, it is required to use a steel plate (steel material) having excellent earthquake resistance as a material in a building structure or the like. In addition, research results so far have shown that steel plates with lower yield ratios have better earthquake resistance.

このため、建築構造物には、降伏比(YR)が80%以下の低降伏比鋼材を使用することが義務付けられている(新耐震設計法1981年施工)。さらに、最近では、建築構造物の高層化や大スパン化などに伴い、建築構造物に、従来より高い強度を有する550MPa級高張力鋼材を適用する事例が増加している。   For this reason, it is obliged to use low yield ratio steel with a yield ratio (YR) of 80% or less for building structures (constructed in 1981). Furthermore, recently, with the increase in the height and span of building structures, there are an increasing number of cases in which 550 MPa class high-tensile steel material having higher strength than before is applied to building structures.

従来、低降伏比を有する550MPa級以上の高張力鋼材は、二相域加熱処理や焼戻処理などの熱処理を施して製造されるのが一般的であった。しかし、熱処理を施すことは、工程が複雑となり製造工期が長期化して、製造コストが高騰するという問題を残していた。このため、上記した二相域加熱処理や焼戻処理を省略した非調質低降伏比高張力鋼材の検討が進められてきた。   Conventionally, a high-tensile steel material of 550 MPa class or higher having a low yield ratio is generally manufactured by performing a heat treatment such as a two-phase region heat treatment or a tempering treatment. However, the heat treatment has a problem in that the process becomes complicated, the manufacturing period is prolonged, and the manufacturing cost is increased. For this reason, the examination of the non-tempered low yield ratio high-tensile steel material in which the above-described two-phase region heat treatment and tempering treatment are omitted has been advanced.

例えば、特許文献1には、C:0.02〜0.04%、固溶B:0.0002〜0.002%を含有し、合金元素含有量に関係する式CENが0.21〜0.30%の範囲を満足する組成と、ベイナイトを主体とし、島状マルテンサイトを0.8〜2.5体積%分散させた組織からなる590MPa級の非調質型低降伏比高張力鋼板が提案されている。特許文献1に記載された技術では、制御圧延のみで製造するとしている。しかし、特許文献1に記載された技術では、鋼板のC含有量を0.02〜0.04%と低炭素化しており、そのため、所望の高強度を確保するためにさらに合金元素量を多量に含有する必要があり、製造コストの高騰を招くという問題がある。   For example, Patent Document 1 contains C: 0.02 to 0.04%, solute B: 0.0002 to 0.002%, and the formula CEN related to the alloy element content is 0.21 to 0. A 590 MPa class non-tempered low yield ratio high tensile strength steel sheet composed of a composition satisfying a range of 30% and a structure mainly composed of bainite and dispersed in an amount of 0.8 to 2.5% by volume of island martensite. Proposed. In the technique described in Patent Document 1, the production is performed only by controlled rolling. However, in the technique described in Patent Document 1, the carbon content of the steel sheet is reduced to 0.02 to 0.04%, so that a large amount of alloying element is required to ensure the desired high strength. Therefore, there is a problem that the manufacturing cost increases.

また、特許文献2には、C:0.045〜0.08%、Si:0.05〜0.50%、Mn:0.6〜2.0%を含み、P、S、Al、Nを調整して含有し、さらにMo及び/又はWを特定の関係式を満足するように含有し、Pcmが0.22%以下となる組成と、板厚中央部の組織が、フェライトを主相とし、20体積%以下の、島状マルテンサイト(MA相)を主とする硬質相を含む複合組織である低降伏比を有する高張力厚鋼板が記載されている。このような組織とすることにより、所望の低降伏比が実現できるとしている。また、このような組織とするために、特許文献2に記載された技術では、上記した組成の鋼素材を、圧延終了温度を表面温度で800〜950℃とする熱間圧延と、0.5〜50℃/sの平均冷却速度で580〜670℃の温度範囲まで加速冷却する冷却処理とを順次施すことが好ましいとしている。しかし、特許文献2に記載された技術では、高価なMo、Wを含有させることを必要とし、製造コストの高騰を招くという問題がある。   Patent Document 2 includes C: 0.045 to 0.08%, Si: 0.05 to 0.50%, Mn: 0.6 to 2.0%, P, S, Al, N In addition, a composition containing Mo and / or W so as to satisfy a specific relational expression, a composition in which Pcm is 0.22% or less, and a structure in the central portion of the plate thickness is the main phase of ferrite. And a high-tensile thick steel plate having a low yield ratio, which is a composite structure containing a hard phase mainly composed of island-shaped martensite (MA phase) of 20% by volume or less. With such a structure, a desired low yield ratio can be realized. Further, in order to obtain such a structure, in the technique described in Patent Document 2, the steel material having the above composition is subjected to hot rolling in which the rolling end temperature is 800 to 950 ° C. at the surface temperature, and 0.5 It is said that it is preferable to sequentially perform a cooling process of accelerated cooling to a temperature range of 580 to 670 ° C. at an average cooling rate of ˜50 ° C./s. However, the technique described in Patent Document 2 has a problem that it requires expensive Mo and W to be contained, resulting in an increase in manufacturing cost.

特許文献3には、C:0.03〜0.30%、Si:0.05〜0.60%、Mn:0.50〜2.5%、Al:0.005〜0.1%を含む鋼を、加熱し、圧延終了温度を900℃〜Ar3変態点の範囲の温度とし該温度域での累積圧下率を30%未満とする熱間圧延と、熱間圧延後空冷し、表面温度が(Ar3変態点−20℃)〜(Ar3変態点−80℃)の範囲の温度となってから水冷を開始し350〜600℃間で冷却停止する加速冷却を施す、低降伏比非調質鋼の製造方法が記載されている。   Patent Document 3 includes C: 0.03 to 0.30%, Si: 0.05 to 0.60%, Mn: 0.50 to 2.5%, Al: 0.005 to 0.1%. The steel containing is heated, the rolling end temperature is set to a temperature in the range of 900 ° C. to Ar 3 transformation point, and the cumulative rolling reduction in the temperature range is set to less than 30%. Is subjected to accelerated cooling in which water cooling is started and cooling is stopped between 350 to 600 ° C. after reaching a temperature in the range of (Ar 3 transformation point −20 ° C.) to (Ar 3 transformation point −80 ° C.). A method for manufacturing steel is described.

特許文献4には、C:0.03〜0.30%、Si:0.05〜0.60%、Mn:0.50〜2.5%、Al:0.005〜0.1%を含む鋼を、加熱し、圧延終了温度を900℃〜Ar3変態点の範囲の温度とし、該温度域での累積圧下率を30%未満とする熱間圧延と、熱間圧延後空冷し、表面温度が(Ar3変態点−20℃)〜(Ar3変態点−80℃)の範囲の温度となってから水冷を開始し250℃以下になるまで加速冷却を施し、その後焼戻し熱処理を行う、低降伏比非調質鋼の製造方法が記載されている。   Patent Document 4 includes C: 0.03 to 0.30%, Si: 0.05 to 0.60%, Mn: 0.50 to 2.5%, Al: 0.005 to 0.1%. The steel containing is heated, the rolling end temperature is set to a temperature in the range of 900 ° C. to Ar 3 transformation point, the hot rolling in which the cumulative rolling reduction in the temperature range is less than 30%, and air cooling after hot rolling, the surface Low yield after the temperature is in the range of (Ar3 transformation point-20 ° C) to (Ar3 transformation point-80 ° C), water cooling is started and accelerated cooling is performed until the temperature falls to 250 ° C or lower, followed by tempering heat treatment. A method for producing non-tempered steel is described.

特許文献5には、C:0.01〜0.20%、Si:0.6%以下、Mn:0.50〜2.2%、Al:0.001〜0.1%、Nb:0.003〜0.030%、Ti:0.005〜0.020%、N:0.006%以下を含む鋼片を、900℃以下の累積圧下量が30%以上で仕上温度がAr3+100℃以下Ar3以上となる熱間圧延を行い、鋼板を(Ar3−20℃)〜(Ar3−100℃)まで空冷し、この温度から水冷を開始し、400〜550℃の範囲で冷却を停止する、低降伏比非調質鋼の製造方法が記載されている。   In Patent Document 5, C: 0.01 to 0.20%, Si: 0.6% or less, Mn: 0.50 to 2.2%, Al: 0.001 to 0.1%, Nb: 0 0.003 to 0.030%, Ti: 0.005 to 0.020%, N: 0.006% or less of steel slabs, the cumulative reduction amount of 900 ° C or less is 30% or more and the finishing temperature is Ar3 + 100 ° C or less Perform hot rolling to be Ar3 or higher, air-cool the steel sheet to (Ar3-20 ° C) to (Ar3-100 ° C), start water cooling from this temperature, stop cooling in the range of 400-550 ° C, low A method for producing a yield ratio non-tempered steel is described.

特許文献3〜5に記載された技術では、合金元素添加量を削減するために、加速冷却を活用して高強度化を図り、高強度と低降伏比を両立させている。これらの技術では、鋼片にAr3変態点以上で圧延を完了する熱間圧延を施した後、加速冷却を開始する前に、オーステナイト+フェライトの二相域温度まで空冷して初析フェライトを生成させることによって低降伏比化を図っている。しかし、これらの技術では、空冷中に生成する初析フェライトと硬質第2相の微細化を図るのが難しく、特に初析フェライト生成量の多い表層部の靱性が低下しやすいという問題があった。また、僅かな冷却開始温度の違いによっても、フェライト生成率が異なってくるため、鋼板ごとの材質ばらつきが大きくなり、安定した鋼板製造が難しいという問題があった。   In the techniques described in Patent Documents 3 to 5, in order to reduce the addition amount of alloy elements, high strength is achieved by utilizing accelerated cooling, and both high strength and low yield ratio are achieved. In these technologies, after hot rolling is completed on the steel slab at the Ar3 transformation point or higher, before starting accelerated cooling, air cooling to the austenite + ferrite two-phase temperature range generates proeutectoid ferrite This is intended to reduce the yield ratio. However, in these techniques, it is difficult to refine the pro-eutectoid ferrite and hard second phase that are generated during air cooling, and there is a problem that the toughness of the surface layer portion where a large amount of pro-eutectoid ferrite is generated tends to decrease. . In addition, since the ferrite generation rate varies depending on a slight difference in cooling start temperature, there is a problem that material variations vary from steel plate to steel plate, making it difficult to produce a stable steel plate.

また、特許文献6には、C:0.01〜0.20%、Si:0.01〜1.0%、Mn:0.1〜2.0%、Al:0.001〜0.1%、N:0.001〜0.010%を含む鋼片に、加熱し900℃までの範囲で累積圧下率が10〜80%の粗圧延と、粗圧延後、2〜40℃/sの加速冷却を(Ar3変態点+50℃)〜(Ar3変態点−50℃)まで行いオーステナイト(γ)相を過冷し、さらに累積圧下率が30〜90%の仕上圧延を650℃以上で終了し、さらに、5〜40℃/sの加速冷却を250〜450℃まで行う低降伏比高張力鋼材の製造方法が記載されている。   In Patent Document 6, C: 0.01 to 0.20%, Si: 0.01 to 1.0%, Mn: 0.1 to 2.0%, Al: 0.001 to 0.1 %, N: In a steel slab containing 0.001 to 0.010%, rough rolling with a cumulative rolling reduction of 10 to 80% in the range up to 900 ° C., and after rough rolling, 2-40 ° C./s Accelerated cooling is performed from (Ar3 transformation point + 50 ° C.) to (Ar3 transformation point−50 ° C.) to supercool the austenite (γ) phase, and finish rolling with a cumulative rolling reduction of 30 to 90% is completed at 650 ° C. or more. Furthermore, a method for producing a low yield ratio high-tensile steel material that performs accelerated cooling at 5 to 40 ° C./s up to 250 to 450 ° C. is described.

特許文献6に記載された技術では、粗圧延の後に加速冷却を行って、γ相をAr3温度付近まで過冷却したうえで、仕上圧延を行うことにより、過冷されたγ相から微細なフェライト(α)を生成させ、さらに仕上圧延後に加速冷却を行うことで、軟質相であるフェライト(α)の微細化と、軟質相と硬質相の比率を適切に制御して高靭性と低降伏比化を両立させることができるとともに、生産性の向上が可能となるとしている。この技術によれば、高価な合金元素の多量含有や生産性の低い複雑な熱処理を必要とすることなく、低降伏比高張力鋼材が製造できるとしている。   In the technique described in Patent Document 6, accelerated cooling is performed after rough rolling, and the γ phase is subcooled to the vicinity of the Ar3 temperature, and then finish rolling is performed, whereby fine ferrite is formed from the supercooled γ phase. (Α) is generated, and accelerated cooling is performed after finish rolling, so that the ferrite (α), which is a soft phase, is refined, and the ratio of the soft phase to the hard phase is appropriately controlled to achieve high toughness and a low yield ratio. It is said that it will be possible to improve the productivity and at the same time. According to this technology, it is said that a high yield steel material with a low yield ratio can be produced without requiring a large amount of expensive alloy elements or complicated heat treatment with low productivity.

また、特許文献7には、Ac3点以上の温度の鋼片または鋼板を、表層から少なくとも板厚方向に製品時板厚の1mm〜30%の領域(表層部)を2℃/s以上の冷却速度でAr1点以下まで急冷し、該表層部がAr3点以下の温度になってから圧延を開始若しくは再開し、(Ac3−50℃)〜Ac3の範囲で圧延を終了し、その後Ac3点以上に復熱することなく、当該表層部をAr1点まで1℃/s以上で冷却し、さらに(Ac1−100℃)〜Ac1の範囲で3min以上滞留させる表層低降伏強度鋼板の製造方法が記載されている。これにより、鋼板板厚の1mm〜30%までの表裏層部の組織が、板厚内部のフェライト粒径の3倍以上の粒径を有するものとなり、降伏強さが板厚内層の降伏強さより5kg/mm以上低く、表層低降伏強度鋼板となるとしている。 Further, in Patent Document 7, a steel piece or steel plate having a temperature of Ac3 or higher is cooled at least 2 ° C./s in a region (surface layer portion) of 1 mm to 30% of the product thickness in the thickness direction from the surface layer. Rapid cooling to the Ar1 point or less at a speed, rolling starts or resumes after the surface layer reaches a temperature of the Ar3 point or less, finishes rolling in the range of (Ac3-50 ° C.) to Ac3, and then reaches the Ac3 point or more. A method for producing a surface low yield strength steel sheet is described in which the surface layer is cooled to 1 ° C./s or more to the Ar1 point without reheating, and further retained for 3 minutes or more in the range of (Ac1-100 ° C.) to Ac1. Yes. As a result, the structure of the front and back layer portions of 1 mm to 30% of the steel plate thickness has a grain size that is three times or more the ferrite grain size inside the plate thickness, and the yield strength is greater than the yield strength of the inner thickness layer. It is said that the surface layer has a low yield strength steel plate having a lower yield of 5 kg / mm 2 or more.

特開2000−219934号公報Japanese Patent Laid-Open No. 2000-219934 特開2007−177325号公報JP 2007-177325 A 特開昭63−219523号公報JP-A-63-219523 特開昭63−223123号公報JP 63-223123 A 特開平1−301819号公報JP-A-1-301818 特開平10−306316号公報JP-A-10-306316 特開平6−49596号公報Japanese Unexamined Patent Publication No. 6-49596

ところで、建築構造物では、柱−梁接合部や柱−ダイアフラム接合部などが多数存在し、多数のT継手や十字継手が形成されている。このようなT継手部や十字継手部では、地震による揺れで変形が生じた時に、溶接止端部など鋼板表面に大きな歪が集中する。   By the way, in a building structure, there are many column-beam joints, column-diaphragm joints, etc., and many T joints and cross joints are formed. In such a T joint part and a cross joint part, when a deformation | transformation arises by the shake by an earthquake, a big distortion concentrates on steel plate surfaces, such as a weld toe part.

図1に、地震による引張・圧縮繰り返し変形を受けた場合に、プレスコラム(冷間成形角形鋼管)や円形鋼管を用いた柱1と通しダイアフラム2の接合部(十字継手)が破壊する状況を模式的に示す。接合部が引張・圧縮繰り返し変形を受けると、通常、溶接部3の溶接止端部で延性亀裂が発生し、該延性亀裂が柱1の板厚中央に向かって伝播(進展)して最終破断に至る。なお、符号4は裏当金である。   Fig. 1 shows the situation where a joint (cross joint) between a column 1 and a diaphragm 2 using a press column (cold-formed square steel pipe) or a circular steel pipe breaks when subjected to repeated tensile and compression deformation due to an earthquake. This is shown schematically. When the joint is subjected to repeated tensile and compression deformations, a ductile crack usually occurs at the weld toe of the weld 3, and the ductile crack propagates (proliferates) toward the thickness center of the column 1 and finally breaks. To. Reference numeral 4 is a backing gold.

このため、破断に至るまでの変形量を大きくするには、柱の表層付近の材質、すなわち、鋼板表層部の延性・靭性が優れていることが重要となるが、建築構造物で使用される冷間曲げ加工によって成形された円形鋼管やプレスコラムの場合、冷間曲げ加工によって鋼板表層付近が著しく硬化し、鋼板を無加工のまま使用する場合と比べて、表層付近の延性・靭性が低下した状態となっている。   For this reason, in order to increase the amount of deformation until breakage, it is important that the material in the vicinity of the surface layer of the column, that is, the ductility and toughness of the steel plate surface layer portion is excellent, but it is used in building structures. In the case of circular steel pipes and press columns formed by cold bending, the steel plate surface area is significantly hardened by the cold bending process, and the ductility and toughness near the surface layer are reduced compared to the case where the steel plate is used without being processed. It has become a state.

そのため、冷間成形角形鋼管を柱に用いた設計は「冷間成形角形鋼管設計・施工マニュアル」(建築センター)によって運用され、2009年には「2008年版冷間成形角形鋼管設計・施工マニュアル」(建築センター)に準拠した設計法として、冷間成形角形鋼管を柱に使用する場合の構造規程が告示されている。   Therefore, the design using a cold-formed square steel pipe as a column is operated by the “Cold-formed square steel pipe design and construction manual” (Architecture Center). In 2009, the “2008 cold-formed square steel pipe design and construction manual” As a design method compliant with (Building Center), structural regulations for using cold-formed square steel pipes for columns are announced.

しかしながら、最近、建築構造物は安全性の他にデザイン性も重視されるようになったため、上記構造規程を緩和させうる、冷間加工を施された後においても、表層付近の延性・靭性の低下が少ない低降伏比高張力厚鋼板を用いた円形鋼管やプレスコラムが要望されている。   However, recently, building structures have become more important in design than safety, so ductility and toughness in the vicinity of the surface layer can be relaxed even after cold working, which can relax the above structural regulations. There is a demand for circular steel pipes and press columns using low yield ratio, high-tensile thick steel plates with little reduction.

特許文献1〜6に記載された技術は、いずれも、全厚引張試験片または板厚1/4tや1/2t位置での丸棒引張試験片により評価される機械的特性(引張特性、延性、靭性)を所望の特性とすることを目的としてなされた技術であり、板厚1/4tや1/2t位置の丸棒引張試験片による引張特性は鋼板表層付近での特性を示唆するものではないから、その特性も不明である。   The techniques described in Patent Documents 1 to 6 are all mechanical properties (tensile properties, ductility) evaluated by full thickness tensile test pieces or round bar tensile test pieces at a thickness of 1/4 t or 1/2 t. , Toughness) is a technique aimed at making it a desired property, and the tensile properties by round bar tensile specimens at the thickness of 1 / 4t or 1 / 2t do not suggest properties near the surface layer of the steel plate. Because there is no, its characteristics are also unknown.

また、特許文献1〜6による制御圧延や加速冷却(TMCP技術)による鋼板は例えば図2に示すように、表層の硬さが最も高く、板厚中央の硬さが最も小さくなる板厚方向の硬さ分布を有している。このような板厚方向硬さ分布を有する鋼板に冷間曲げ加工を施すと、表裏面近傍の硬さがさらに増加して、図2の「曲げ加工後」のような硬さ分布となり、板厚中央部と表層部の硬さの差がさらに拡大する。   Moreover, as shown in FIG. 2, for example, as shown in FIG. 2, the steel sheet by controlled rolling or accelerated cooling (TMCP technology) according to Patent Documents 1 to 6 has the highest surface layer hardness and the lowest center thickness. It has a hardness distribution. When cold bending is performed on a steel sheet having such a thickness distribution in the plate thickness direction, the hardness in the vicinity of the front and back surfaces is further increased, resulting in a hardness distribution like “after bending” in FIG. The difference in hardness between the thickness center portion and the surface layer portion is further enlarged.

特許文献7に記載された技術によれば、鋼板の表裏層部を低降伏強さとすることができ、冷間曲げ加工後の鋼板表層部の延性を向上させることができるが、鋼板の表層部のフェライト粒が粗大であるため、靭性が十分であるとはいえず、部材として建築構造物に組み入れられた場合、該部材から脆性破壊を発生させることが懸念される。   According to the technique described in Patent Document 7, the front and back layer portions of the steel sheet can have low yield strength, and the ductility of the steel sheet surface layer portion after cold bending can be improved. Since the ferrite grains are coarse, it cannot be said that the toughness is sufficient, and when incorporated as a member in a building structure, there is a concern that brittle fracture will occur from the member.

そこで、本発明は、上記した従来技術の問題を解決し、焼入焼戻や焼準等の熱処理を施すことなく、また合金含有量を最小限に抑制したうえで、プレスコラムや円形鋼管を用いた建築構造物部材用として好適な、冷間曲げ加工後においても鋼板表層部の硬さ増加が少なく、鋼板表層部の延性、靭性に優れた、降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:75%以下を有する非調質低降伏比高張力厚鋼板およびその製造方法を提供することを目的とする。   Therefore, the present invention solves the above-mentioned problems of the prior art, does not perform heat treatment such as quenching and tempering, and normalization, and suppresses the alloy content to a minimum, and then press columns and circular steel pipes. Suitable for the building structure member used, even after cold bending, there is little increase in the hardness of the steel sheet surface layer part, and the steel sheet surface layer part is excellent in ductility and toughness, yield strength: 385 MPa or more, tensile strength: An object of the present invention is to provide a non-tempered low yield ratio high tensile steel plate having a yield ratio of 550 MPa or more and a yield ratio of 75% or less, and a method for producing the same.

本発明者らは、上記目的を達成するため、板厚方向に不均一な硬さ分布をある程度許容することを前提とし、方針1.冷間曲げ加工による塑性歪は、鋼板の表裏面で最大、板厚中央付近の中立点ではゼロとなり、冷間曲げによる加工硬化は鋼板表層部で最も顕著となるため、冷間曲げ加工前の板厚方向硬さ分布を制御し、まず、表層部付近の硬さを低下する。方針2.板厚中央部の硬さをそのままにして表層部の硬さを低下すれば、鋼板全厚での強度が低下してしまうので、板厚中央部で一定以上の硬さ(強度)を確保する。以上の方針により、鋭意検討を行い、以下の知見を得た。
1.表層部付近の硬さを低下させ、さらには低降伏比を達成するためには、少なくとも表層部のミクロ組織において、軟質相であるフェライト(好ましくは50面積%以上)を析出させ、硬質相との複相組織とすることが必要で、さらに、表層部におけるフェライトの平均粒径を所望の適正範囲内に調整することにより、表層部の延性・靭性を所望の範囲内とすることができる。
2.下記(1)〜(4)を満足する厚鋼板であれば、冷間曲げ加工後にも建築構造物部材用として必要な変形性能を確保できる。
(1)鋼板の、少なくとも表層部(表面および裏面から板厚方向に1〜5mmの領域)ミクロ組織をフェライトおよび硬質相からなる複相組織とすること。
(2)鋼板表層部の平均硬さが225HV以下を満足すること。
(3)鋼板表層部と板厚中央部の硬度差が60HV以下であること。
(4)鋼板表層部の平均フェライト粒径が4.0〜18.0μmの範囲を満足すること。
ここで、硬質相とはパーライト、ベイナイト、マルテンサイトのうちの1種または2種以上からなる相を意味し、鋼板表層部とは鋼板表裏面から板厚方向に1〜5mmの領域を、また、板厚中央部とは板厚中心±2mmの領域を指す。なお、鋼板表層部の組織、硬さを限定した理由は、溶接構造物の破壊に対しては、鋼板表面または裏面から板厚方向に1〜5mmの領域である表層部の影響が大きいことを見出したことに基づく。
In order to achieve the above object, the present inventors assume that a non-uniform hardness distribution in the thickness direction is allowed to some extent. The plastic strain due to cold bending is maximum at the front and back surfaces of the steel sheet, zero at the neutral point near the center of the plate thickness, and work hardening by cold bending is most noticeable at the surface layer of the steel sheet. The thickness distribution in the thickness direction is controlled, and first, the hardness near the surface layer is reduced. Policy 2. If the hardness of the surface layer is reduced while leaving the hardness at the center of the plate thickness as it is, the strength at the full thickness of the steel plate will decrease, so ensure a certain level of hardness (strength) at the center of the plate thickness. . Based on the above policy, intensive studies were conducted and the following findings were obtained.
1. In order to reduce the hardness in the vicinity of the surface layer portion and achieve a low yield ratio, at least in the microstructure of the surface layer portion, ferrite (preferably 50 area% or more) as a soft phase is precipitated, In addition, the ductility and toughness of the surface layer portion can be set within the desired range by adjusting the average grain size of ferrite in the surface layer portion within a desired appropriate range.
2. If it is a thick steel plate which satisfies the following (1)-(4), the deformation | transformation performance required as an object for building structure members can be secured even after cold bending.
(1) At least the surface layer portion (region of 1 to 5 mm in the plate thickness direction from the front surface and the back surface) of the steel sheet has a multiphase structure composed of ferrite and a hard phase.
(2) The average hardness of the steel sheet surface layer portion satisfies 225 HV or less.
(3) The hardness difference between the steel plate surface layer and the plate thickness center is 60HV or less.
(4) The average ferrite grain size of the steel sheet surface layer portion satisfies the range of 4.0 to 18.0 μm.
Here, the hard phase means a phase composed of one or more of pearlite, bainite, and martensite, and the steel sheet surface layer portion refers to a region of 1 to 5 mm in the sheet thickness direction from the steel sheet front and back surfaces. The plate thickness central portion refers to a plate thickness center of ± 2 mm. In addition, the reason for limiting the structure and hardness of the steel sheet surface layer part is that the influence of the surface layer part, which is an area of 1 to 5 mm in the thickness direction from the steel sheet front surface or the back surface, is great for the destruction of the welded structure. Based on finding.

鋼板表面または裏面から板厚方向に1mm未満の領域である最表層を除外したのは、最表層が、圧延や加速冷却などによって極めて複雑な熱履歴を受けるため、最表層部のミクロ組織を制御することは極めて困難な場合が多いためである。   Excluding the outermost layer, which is less than 1 mm in the thickness direction from the front or back surface of the steel sheet, the outermost layer receives an extremely complicated thermal history due to rolling, accelerated cooling, etc., so the microstructure of the outermost layer is controlled. This is because it is often extremely difficult to do.

さらに、本発明者らは、冷間加工後の表層部延性・靭性の向上に加えて、大入熱溶接熱影響部の靭性向上に及ぼす各種要因ついて鋭意研究した。その結果、Nb、Moの含有が、大入熱溶接熱影響部の靭性を著しく劣化させることを見出した。Nb、Moは、焼入れ性を向上させる元素であり、島状マルテンサイトを含む、上部ベイナイトの生成に大きく寄与し、大入熱溶接熱影響部の靭性を著しく劣化させる。そこで、大入熱溶接熱影響部の靭性向上のために、本発明では、Nb、Moを添加することなく、さらに不純物としてもNb、Moの含有を厳しく制限することが必要であるという知見を得た。   Furthermore, the present inventors diligently studied various factors affecting the improvement of the toughness of the high heat input welding heat affected zone in addition to the improvement of the surface layer ductility and toughness after cold working. As a result, it has been found that the inclusion of Nb and Mo significantly deteriorates the toughness of the high heat input welding heat-affected zone. Nb and Mo are elements that improve the hardenability, greatly contribute to the formation of upper bainite including island martensite, and significantly deteriorate the toughness of the heat-affected zone with high heat input welding. Therefore, in order to improve the toughness of the high heat input welding heat-affected zone, in the present invention, it is necessary to strictly limit the content of Nb and Mo as impurities without adding Nb and Mo. Obtained.

本発明は、かかる知見に基づいて、さらに検討を加えて完成されたもので、すなわち、本発明は、
1.質量%で、C:0.05〜0.16%、Si:0.05〜0.45%、Mn:1.2〜1.8%、P:0.020%以下、S:0.005%以下、Al:0.05%以下、Ti:0.005〜0.020%、N:0.0040%以下、4.0≧Ti/N≧2.0、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%未満に制限し、さらに下記(1)式で定義される炭素当量Ceqが、0.35〜0.48を満足し、残部Feおよび不可避的不純物からなる成分組成と、少なくとも鋼板の表層部において、平均結晶粒径が4.0〜18.0μmフェライトと、パーライト、ベイナイトおよびマルテンサイトの1種または2種以上からなる硬質相からなり、フェライト面積率が45%〜70%のミクロ組織を有し、鋼板の表層部の平均硬さが225HV以下で、該表層部と板厚中央部との硬度差が60HV以下であることを特徴とする、冷間加工後の表層部の延性・靭性に優れる降伏強さ385MPa以上、引張強さ550MPa以上、降伏比75%以下である非調質低降伏比高張力厚鋼板。
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 (1)
(ここで、C、Mn、Cr、Mo、V、Cu、Niは各元素の含有量(質量%)で含有しない場合は0とする。)
2.成分組成が、さらに質量%で、Cu:0.05〜0.50%、Ni:0.05〜0.80%、Cr:0.05〜0.60%、V:0.01〜0.05%、B:0.0003〜0.0030%の1種または2種以上を含有することを特徴とする1に記載の非調質低降伏比高張力厚鋼板。
3.成分組成が、更に質量%で、Ca:0.0005〜0.0050%、REM:0.0010〜0.0050%の1種または2種を含有することを特徴とする1または2に記載の非調質低降伏比高張力厚鋼板。
4.更に、下記(2)式で定義されるACRが0.2〜0.8であることを特徴とする3に記載の非調質低降伏比高張力厚鋼板。
ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S) (2)
(ここで、Ca、REM、O、Sは各元素の含有量(質量%))
5.成分組成が1乃至4のいずれか一つに記載の鋼素材を1050〜1200℃に加熱後、表面温度で950℃以下の温度域での累積圧下量が30%以上で、圧延終了温度が表面温度で900℃以下Ar3変態点以上となる熱間圧延を行い、その後、第一段冷却とし表面温度でAr3変態点以上の温度から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、表面温度が(Ar3変態点−100℃)以下550℃以上となるまで加速冷却し、冷却停止後復熱させ、第二段冷却として表面温度が(Ar3変態点−20℃)以下600℃以上、かつ、表面温度が極大値をとった時点から、式(3)を満たす時間t1(秒)以上、式(4)を満たす時間t2(秒)以内から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止後の復熱で表面温度が600℃以下400℃以上になる冷却停止温度まで加速冷却することを特徴とする非調質低降伏比高張力厚鋼板の製造方法。
The present invention has been completed with further studies based on this finding, that is, the present invention
1. In mass%, C: 0.05 to 0.16%, Si: 0.05 to 0.45%, Mn: 1.2 to 1.8%, P: 0.020% or less, S: 0.005 %: Al: 0.05% or less, Ti: 0.005-0.020%, N: 0.0040% or less, 4.0 ≧ Ti / N ≧ 2.0, and Nb, Mo as impurity elements Nb: 0.004% or less and Mo: less than 0.04%, and further, the carbon equivalent Ceq defined by the following formula (1) satisfies 0.35 to 0.48, and the balance Fe and Consists of component composition consisting of unavoidable impurities, and at least in the surface layer portion of the steel sheet, the ferrite has an average crystal grain size of 4.0 to 18.0 μm, and a hard phase composed of one or more of pearlite, bainite and martensite. Has a microstructure with a ferrite area ratio of 45% to 70% The average hardness of the surface layer portion of the steel sheet is 225 HV or less, and the difference in hardness between the surface layer portion and the plate thickness central portion is 60 HV or less, which is excellent in the ductility and toughness of the surface layer portion after cold working A non-tempered low yield ratio high tensile steel plate having a yield strength of 385 MPa or more, a tensile strength of 550 MPa or more, and a yield ratio of 75% or less.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
(Here, C, Mn, Cr, Mo, V, Cu, and Ni are set to 0 when not contained in the content (mass%) of each element.)
2. The component composition is further mass%, Cu: 0.05-0.50%, Ni: 0.05-0.80%, Cr: 0.05-0.60%, V: 0.01-0. The non-tempered low yield ratio high-tensile thick steel plate according to 1, which contains one or more of 05% and B: 0.0003 to 0.0030%.
3. 3. The component composition according to 1 or 2, further comprising, by mass%, one or two of Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% Non-tempered, low yield ratio, high tensile thick steel plate.
4). Furthermore, ACR defined by the following formula (2) is 0.2 to 0.8, the non-tempered low yield ratio high-tensile thick steel plate according to 3.
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
(Ca, REM, O, and S are the contents of each element (mass%))
5. After heating the steel material according to any one of the components 1 to 4 to 1050 to 1200 ° C., the cumulative reduction amount in the temperature range of 950 ° C. or less at the surface temperature is 30% or more, and the rolling end temperature is the surface The hot rolling is performed at a temperature of 900 ° C. or lower and the Ar3 transformation point or higher, and then the first stage cooling is performed, and the average cooling rate 2 at the 1/4 position of the sheet thickness (t) from the temperature of the Ar3 transformation point or higher at the surface temperature. Acceleration cooling until the surface temperature becomes 550 ° C. or higher at a temperature of not less than ℃ / s and (Ar3 transformation point−100 ° C.). ) From 600 ° C. or higher and when the surface temperature reaches the maximum value, the thickness (t) from the time t1 (second) that satisfies the formula (3) to the time t2 (second) that satisfies the formula (4) ) At an average cooling rate at 1/4 position of 2 ° C./s or higher, after cooling stop Non-heat treated low yield ratio method for manufacturing a high-tensile steel plates, characterized by accelerated cooling to a cooling stop temperature that the surface temperature by heat is 600 ° C. or less 400 ° C. or higher.

6.前記第一段冷却は、表面温度でAr3変態点以上の温度から冷却を開始し、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止温度が表面温度で550℃以上となる加速冷却を、複数回繰り返す冷却とし、該複数回の加速冷却において、冷却停止温度が表面温度で(Ar3変態点−100℃)以下550℃以上となる加速冷却を少なくとも1回含むことを特徴とする5に記載の非調質低降伏比高張力厚鋼板の製造方法。
7.第二段冷却を、表面温度が(Ar3変態点−20℃)以下600℃以上、かつ、表面温度が極大値をとった時点から、式(3)を満たす時間t1(秒)以上、式(4)を満たす時間t2(秒)以内から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止後復熱で表面温度が400℃以上となる冷却停止温度まで加速冷却する冷却を、複数回繰り返す冷却とし、前記複数回の加速冷却において、冷却停止後の復熱で表面温度が600℃以下400℃以上になる冷却停止温度まで冷却する加速冷却を最終冷却とすることを特徴とする5または6に記載の非調質低降伏比高張力厚鋼板の製造方法。
6). The first stage cooling starts from a temperature equal to or higher than the Ar3 transformation point at the surface temperature, is at an average cooling rate of 2 ° C./s or more at a 1/4 position of the plate thickness (t), and the cooling stop temperature is the surface temperature. The accelerated cooling at 550 ° C. or higher is repeated multiple times, and in the multiple times of accelerated cooling, the accelerated cooling at which the cooling stop temperature is not more than (Ar 3 transformation point−100 ° C.) at the surface temperature is 550 ° C. or higher at least once. 5. The method for producing a non-tempered low yield ratio high-tensile thick steel plate according to 5, wherein
7). In the second stage cooling, the surface temperature is (Ar3 transformation point−20 ° C.) or lower and 600 ° C. or higher, and the time when the surface temperature takes the maximum value is equal to or longer than the time t1 (seconds) that satisfies Equation (3). 4) The cooling stop temperature at which the surface temperature becomes 400 ° C. or higher by reheating after cooling stop at an average cooling rate of 2 ° C./s or more at the 1/4 position of the plate thickness (t) from within the time t2 (seconds) that satisfies The cooling to be accelerated cooling to the cooling is repeated a plurality of times, and the accelerated cooling to the cooling stop temperature at which the surface temperature becomes 600 ° C. or lower and 400 ° C. or higher by the recuperation after the cooling is stopped is the final cooling. The method for producing a non-tempered, low yield ratio, high-tensile thick steel plate according to 5 or 6, wherein

本発明によれば、冷間曲げ加工後においても、鋼板表層部の硬さ増加が少なく、鋼板表層部の延性、靭性に優れ、建築構造物部材用として好適な、降伏強さ:385MPa以上、引張強さ:550MPa以上の高強度と降伏比:75%以下の低降伏比を有する非調質低降伏比高張力厚鋼板を、熱処理を施すことなく、また多量な合金含有を行うことなく、製造でき、産業上格段の効果を奏する。また、本発明になる非調質低降伏比高張力厚鋼板は、鋼構造物の軽量化や、鋼構造物の耐震性の向上に大きく寄与するという効果もある。   According to the present invention, even after cold bending, there is little increase in hardness of the steel sheet surface layer part, excellent ductility and toughness of the steel sheet surface layer part, suitable for building structure members, yield strength: 385 MPa or more, Tensile strength: high strength of 550 MPa or higher and yield ratio: non-refined low yield ratio high tensile steel plate having a low yield ratio of 75% or less, without heat treatment, and without containing a large amount of alloy, It can be manufactured and has a remarkable industrial effect. Moreover, the non-tempered low yield ratio high-tensile thick steel plate according to the present invention also has an effect of greatly contributing to the weight reduction of the steel structure and the improvement of the earthquake resistance of the steel structure.

柱−ダイアフラム接合部における破壊状況を説明する模式図。The schematic diagram explaining the destruction condition in a column-diaphragm junction part. 非調質厚鋼板の板厚方向硬さ分布を説明する模式図。The schematic diagram explaining the thickness direction hardness distribution of a non-tempered thick steel plate. プレスコラム−ダイアフラム接合部の三点曲げ試験(コラム曲げ試験)方法の概略を説明する模式図。The schematic diagram explaining the outline of the three-point bending test (column bending test) method of a press column-diaphragm junction part. 図3の試験体における溶接部近傍を説明する模式図。The schematic diagram explaining the welding part vicinity in the test body of FIG. コラム曲げ試験の要領を説明する模式図。The schematic diagram explaining the point of a column bending test. コラム曲げ試験における荷重−変形量ヒステリシス曲線を説明する模式図。The schematic diagram explaining the load-deformation amount hysteresis curve in a column bending test. 本発明における冷却工程の冷却条件を説明する模式図。The schematic diagram explaining the cooling conditions of the cooling process in this invention.

本発明では、成分組成、ミクロ組織、板厚方向の硬さ分布を規定する。
[成分組成]以下の説明において、%は質量%とする。
C:0.05〜0.16%
Cは、鋼の強度を増加させ、構造用鋼材として必要な強度を確保するのに有用な元素である。さらにCは、硬質相の体積率を増加させ、降伏比を低下させる作用を有する。このような効果を得るためには0.05%以上の含有を必要とする。一方、0.16%を超える含有は、溶接性と靭性を顕著に低下させる。このため、Cは0.05〜0.16%の範囲に限定した。なお、好ましくは0.06〜0.15%である。
In the present invention, the component composition, microstructure, and hardness distribution in the thickness direction are defined.
[Component Composition] In the following description, “%” means “mass%”.
C: 0.05 to 0.16%
C is an element useful for increasing the strength of steel and ensuring the strength required as a structural steel material. Further, C has an effect of increasing the volume fraction of the hard phase and decreasing the yield ratio. In order to acquire such an effect, 0.05% or more of content is required. On the other hand, if the content exceeds 0.16%, weldability and toughness are significantly reduced. For this reason, C was limited to the range of 0.05 to 0.16%. In addition, Preferably it is 0.06 to 0.15%.

Si:0.05〜0.45%
Siは、脱酸剤として作用するとともに、鋼中に固溶し鋼材の強度を増加させる。このような効果を得るためには0.05%以上の含有を必要とする。一方、0.45%を超える含有は、母材の靱性を低下させるとともに、溶接熱影響部(HAZとも言う)靱性を顕著に低下させる。このため、Siは0.05〜0.45%の範囲に限定した。なお、好ましくは、0.05〜0.35%である。
Si: 0.05 to 0.45%
Si acts as a deoxidizer and dissolves in the steel to increase the strength of the steel material. In order to acquire such an effect, 0.05% or more of content is required. On the other hand, if the content exceeds 0.45%, the toughness of the base metal is lowered and the toughness of the weld heat affected zone (also referred to as HAZ) is significantly lowered. For this reason, Si was limited to the range of 0.05 to 0.45%. In addition, Preferably, it is 0.05 to 0.35%.

Mn:1.2〜1.8%
Mnは、固溶して鋼の強度を増加させる作用を有する元素で安価であり、高価な他の合金元素の含有を最小限に抑えることを目的の一つとする本発明では、所望の高強度(引張強さ550MPa以上)を確保するために、1.2%以上の含有を必要とする。一方、1.8%を超える含有は、母材の靱性およびHAZ靱性を著しく低下させる。このため、Mnは1.2〜1.8%の範囲に限定した。なお、好ましくは1.2〜1.6%である。
Mn: 1.2 to 1.8%
Mn is an element that has the effect of increasing the strength of steel by solid solution and is inexpensive, and in the present invention, which aims to minimize the content of other expensive alloy elements, the desired high strength In order to ensure (tensile strength of 550 MPa or more), a content of 1.2% or more is required. On the other hand, if the content exceeds 1.8%, the toughness and the HAZ toughness of the base material are significantly reduced. For this reason, Mn was limited to 1.2 to 1.8% of range. In addition, Preferably it is 1.2 to 1.6%.

P:0.020%以下
Pは、鋼の強度を増加させる作用を有する元素であるが、靱性、とくに溶接部の靱性を低下させる元素であり、本発明ではできるだけ低減することが望ましいが、過度の低減は、精錬コストを高騰させ経済的に不利となるため、0.005%程度以上とすることが好ましい。一方、0.020%を超えて含有すると、上記した悪影響が顕著となるため、Pは0.020%以下に限定した。なお、好ましくは0.015%以下である。
P: 0.020% or less P is an element that has an effect of increasing the strength of steel, but is an element that lowers toughness, particularly the toughness of a welded portion. Since the reduction of the cost increases the refining cost and is economically disadvantageous, it is preferable to be about 0.005% or more. On the other hand, if the content exceeds 0.020%, the above-described adverse effects become remarkable, so P is limited to 0.020% or less. In addition, Preferably it is 0.015% or less.

S:0.005%以下
Sは、鋼中ではMnS等の硫化物系介在物として存在し、母材および溶接部の靱性を劣化させるとともに、鋳片中央偏析部などに多量に偏在して鋳片等における欠陥を発生しやすくする。このような傾向は0.005%を超える含有で顕著となる。このため、Sは0.005%以下に限定した。好ましくは0.003%以下である。なお、過度のS低減は、精錬コストを高騰させ、経済的に不利となるため、Sは0.001%程度以上とすることが望ましい。
S: 0.005% or less S is present in the steel as sulfide inclusions such as MnS, which deteriorates the toughness of the base metal and the welded portion, and is unevenly distributed in the center segregated portion of the slab. It makes it easier to generate defects in pieces. Such a tendency becomes remarkable when the content exceeds 0.005%. For this reason, S was limited to 0.005% or less. Preferably it is 0.003% or less. In addition, since excessive S reduction raises refining cost and becomes economically disadvantageous, it is desirable for S to be about 0.001% or more.

Al:0.05%以下
Alは、脱酸剤として作用する元素であり、高張力鋼の溶鋼脱酸プロセスにおいては、脱酸剤として、もっとも汎用的に使われる。このような効果を得るためには、0.01%以上含有することが望ましいが、0.05%を超える含有は、母材の靱性が低下するとともに、溶接時に溶接金属に混入して溶接金属部靱性を低下させる。このため、Alは0.05%以下に限定した。なお、好ましくは0.010〜0.045%である。
Al: 0.05% or less Al is an element that acts as a deoxidizer, and is most commonly used as a deoxidizer in a molten steel deoxidation process of high-strength steel. In order to obtain such an effect, it is desirable to contain 0.01% or more. However, if the content exceeds 0.05%, the toughness of the base material decreases, and the weld metal is mixed into the weld metal during welding. Reduce toughness. For this reason, Al was limited to 0.05% or less. In addition, Preferably it is 0.010 to 0.045%.

N:0.0040%以下
Nは、鋼中に固溶している場合には、冷間加工後に歪時効を起こし靭性を劣化させるため、本発明ではできるだけ低減することが望ましい。0.0040%を超えて含有すると、靭性の劣化が著しくなる。このため、Nは0.0040%以下に限定した。
N: 0.0040% or less N, when dissolved in steel, causes strain aging after cold working and deteriorates toughness. Therefore, it is desirable to reduce N as much as possible in the present invention. If the content exceeds 0.0040%, the toughness deteriorates remarkably. For this reason, N was limited to 0.0040% or less.

Ti:0.005〜0.020%
Tiは、Nとの親和力が強い元素であり、凝固時にTiNとして析出し、鋼中の固溶Nを減少させ、冷間加工後のNの歪時効による靭性劣化を低減する作用を有する。また、Tiは、HAZの組織改善を介して、HAZ靭性の向上にも寄与する。このような効果を得るためには、0.005%以上の含有を必要とする。一方、0.020%を超えて含有すると、TiN粒子が粗大化し、上記した効果が期待できなくなる。このため、Tiは0.005〜0.020%の範囲に限定した。なお、好ましくは0.007〜0.015%である。
Ti: 0.005-0.020%
Ti is an element having a strong affinity for N, and precipitates as TiN during solidification, thereby reducing solid solution N in the steel and reducing the toughness deterioration due to strain aging of N after cold working. Ti also contributes to the improvement of HAZ toughness through the improvement of the HAZ structure. In order to acquire such an effect, 0.005% or more of content is required. On the other hand, if the content exceeds 0.020%, the TiN particles become coarse and the above-described effects cannot be expected. For this reason, Ti was limited to 0.005 to 0.020% of range. In addition, Preferably it is 0.007 to 0.015%.

4.0≧Ti/N≧2.0
本発明では、N含有量(質量%)に見合う量のTiを含有させ、固溶NをTiNとして固定する。このため、Ti含有量(質量%)とN含有量(質量%)との比、Ti/Nが2.0以上を満足するように、Ti含有量(質量%)、N含有量(質量%)を調整する。Ti/Nが2.0未満では、N含有量(質量%)に比べてTi含有量(質量%)が少なすぎ、多くのNが固溶Nとして残存して、HAZ靭性が低下、溶接部からの脆性破壊発生により部材変形性能が低下する場合がある。このため、Ti/Nを2.0以上に限定した。一方、Ti/Nが4.0を超えて大きくなると、TiN粒子が粗大化して、所望の効果を確保できなくなる。このため、Ti/Nを2.0〜4.0の範囲に限定した。なお好ましくは、2.5〜3.5の範囲である。
4.0 ≧ Ti / N ≧ 2.0
In the present invention, an amount of Ti commensurate with the N content (mass%) is contained, and the solid solution N is fixed as TiN. Therefore, the ratio of Ti content (mass%) to N content (mass%), Ti content (mass%), N content (mass%) so that Ti / N satisfies 2.0 or more. ). When Ti / N is less than 2.0, the Ti content (mass%) is too small compared to the N content (mass%), so that a lot of N remains as solid solution N, and the HAZ toughness is lowered. Due to the occurrence of brittle fracture, the member deformation performance may deteriorate. For this reason, Ti / N was limited to 2.0 or more. On the other hand, when Ti / N exceeds 4.0, the TiN particles become coarse and the desired effect cannot be ensured. For this reason, Ti / N was limited to the range of 2.0-4.0. In addition, Preferably, it is the range of 2.5-3.5.

Nb:0.004%以下、Mo:0.04%未満
Nb、Moは、焼入れ性を向上する元素であり、島状マルテンサイトを含む上部ベイナイトを生成しやすくして、大入熱溶接熱影響部の靭性を低下させる。このため、本発明では、Nb、Moは添加しない。不可避的不純物として含有される場合は、Nb:0.004%以下、Mo:0.04%未満に限定し、製造コストの許す範囲で少なくする。
Nb: 0.004% or less, Mo: less than 0.04% Nb and Mo are elements that improve the hardenability and facilitate the formation of upper bainite containing island martensite. Reduce the toughness of the part. For this reason, Nb and Mo are not added in the present invention. When it is contained as an inevitable impurity, it is limited to Nb: 0.004% or less and Mo: less than 0.04%, and is reduced within the range allowed by the manufacturing cost.

不可避的不純物として、Nbが0.004%を超えて、または、Moが0.04%を超えて含有すると、大入熱溶接熱影響部の靭性が低下する。なお、Nb:0.004%以下、Mo:0.04%未満を満足させるためには、Nb、Moの含有量が少ない原材料や、溶製炉耐火物等を使用することで肝要である。   If Nb exceeds 0.004% or Mo exceeds 0.04% as an unavoidable impurity, the toughness of the high heat input weld heat affected zone is lowered. In order to satisfy Nb: 0.004% or less and Mo: less than 0.04%, it is important to use raw materials with a low content of Nb and Mo, smelting furnace refractories, and the like.

Ceq:0.35〜0.48
Ceqが、0.35未満では、所望の母材強度を確保できないうえ、溶接熱影響部の軟化を所望の許容限度内に抑えることができない。一方、Ceqが、0.48を超えて高くなると、溶接性が低下するとともに、母材靭性、HAZ靭性が低下する。このため、Ceqは0.35〜0.48の範囲に限定した。なお、好ましくは0.36〜0.46である。なお、Ceqは次式による。
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15(ここで、C、 Mn、Cr、Mo、V、Cu、Niは各元素の含有量(質量%)で含有しない場合は0とする。)
上記した成分が基本成分で残部Feおよび不可避的不純物であるが、さらに、選択元素として、Cu:0.05〜0.50%、Ni:0.05〜0.80%、Cr:0.05〜0.60%、V:0.01〜0.05%、B:0.0003〜0.0030%のうちから選ばれた1種または2種以上、および/または、Ca:0.0005〜0.0050%、REM:0.0010〜0.0050%のうちから選ばれた1種または2種を含有できる。
Ceq: 0.35-0.48
If Ceq is less than 0.35, the desired base metal strength cannot be ensured, and the softening of the weld heat affected zone cannot be suppressed within a desired allowable limit. On the other hand, when Ceq is higher than 0.48, the weldability is lowered and the base material toughness and the HAZ toughness are lowered. For this reason, Ceq was limited to the range of 0.35 to 0.48. In addition, Preferably it is 0.36-0.46. Ceq is based on the following equation.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (where C, Mn, Cr, Mo, V, Cu, Ni are 0 in the case of not containing each element content (mass%)) .)
The above-mentioned components are basic components and the balance is Fe and unavoidable impurities. Further, as selective elements, Cu: 0.05 to 0.50%, Ni: 0.05 to 0.80%, Cr: 0.05 ˜0.60%, V: 0.01 to 0.05%, B: one or more selected from 0.0003 to 0.0030%, and / or Ca: 0.0005 One or two selected from 0.0050% and REM: 0.0010 to 0.0050% can be contained.

Cu:0.05〜0.50%、Ni:0.05〜0.80%、Cr:0.05〜0.60%、V:0.01〜0.05%、B:0.0003〜0.0030%のうちから選ばれた1種または2種以上
Cu、Ni、Cr、Vはいずれも、鋼の強度を増加させる作用を有する元素であり、選択して含有できる。
Cu: 0.05 to 0.50%, Ni: 0.05 to 0.80%, Cr: 0.05 to 0.60%, V: 0.01 to 0.05%, B: 0.0003 to One or more selected from 0.0030% are Cu, Ni, Cr, and V, which are elements having an action of increasing the strength of steel, and can be selected and contained.

Cuは、固溶強化や焼入性向上を介して、鋼板の強度を増加させ、厚鋼板の高強度化に寄与する。このような効果を得るためには、0.05%以上含有することが好ましいが、0.50%を超える含有は、合金コストの増加や熱間脆性による表面性状の劣化を招く.このため、含有する場合には、Cuは0.05〜0.50%の範囲に限定することが好ましい。なお、より好ましくは0.10〜0.40%である。   Cu increases the strength of the steel sheet through solid solution strengthening and hardenability improvement, and contributes to increasing the strength of the thick steel sheet. In order to obtain such an effect, the content is preferably 0.05% or more. However, if the content exceeds 0.50%, the alloy cost increases and surface properties are deteriorated due to hot brittleness. For this reason, when it contains, it is preferable to limit Cu to 0.05 to 0.50% of range. In addition, More preferably, it is 0.10 to 0.40%.

Niは、靱性をほとんど劣化させることなく、鋼板の強度を増加させる元素であり、しかもHAZ靱性への悪影響も小さく、厚鋼板の高強度化に有用な元素である。このような効果を得るためには、0.05%以上含有することが好ましいが、0.80%を超える多量の含有は、Niが高価な元素であるため、合金コストの増加を招く。このため、含有する場合は、Niは0.05〜0.80%に限定することが好ましい。なお、より好ましくは0.10〜0.80%である。   Ni is an element that increases the strength of the steel sheet with little deterioration in toughness, and has a small adverse effect on the HAZ toughness, and is an element useful for increasing the strength of thick steel sheets. In order to acquire such an effect, it is preferable to contain 0.05% or more, but when it contains a large amount exceeding 0.80%, Ni is an expensive element, which causes an increase in alloy cost. For this reason, when it contains, it is preferable to limit Ni to 0.05 to 0.80%. In addition, More preferably, it is 0.10 to 0.80%.

Crは、焼入性向上を介し、母材の強度を増加させる元素であり、厚鋼板の高強度化に有用な元素である。このような効果を得るためには、0.05%以上含有することが好ましいが、0.60%を超える含有は、合金コストの増加を招く。このため、含有する場合には、Crは0.05〜0.60%の範囲に限定することが好ましい。なお、より好ましくは0.10〜0.60%である。   Cr is an element that increases the strength of the base material through improvement in hardenability, and is an element useful for increasing the strength of thick steel plates. In order to acquire such an effect, it is preferable to contain 0.05% or more, but inclusion exceeding 0.60% causes an increase in alloy cost. For this reason, when it contains, it is preferable to limit Cr to 0.05 to 0.60% of range. In addition, More preferably, it is 0.10 to 0.60%.

Vは、析出強化を介して母材の強度を増加させる元素であり、厚鋼板の高強度化のために有用な元素である。このような効果を得るためには、0.01%以上含有することが好ましいが、0.05%を超える含有は、母材やHAZの靭性を低下させる。このため、含有する場合には、Vは0.01〜0.05%の範囲に限定することが好ましい。なお、より好ましくは0.02〜0.04%である。   V is an element that increases the strength of the base metal through precipitation strengthening, and is a useful element for increasing the strength of the thick steel plate. In order to acquire such an effect, it is preferable to contain 0.01% or more, but inclusion exceeding 0.05% reduces the toughness of a base material and HAZ. For this reason, when it contains, it is preferable to limit V to 0.01 to 0.05% of range. In addition, More preferably, it is 0.02-0.04%.

Bは焼入れ性の向上を介し、鋼の強度増加に寄与する元素である。このような効果を得るために、0.0003%以上含有することが好ましいが、0.0030%を超える含有は、母材やHAZ靭性を劣化させる。このため、含有する場合には、Bは0.0003%〜0.0030%の範囲に限定することが好ましい。なお、より好ましくは0.0006〜0.0020%である。   B is an element that contributes to an increase in the strength of steel through the improvement of hardenability. In order to obtain such an effect, the content is preferably 0.0003% or more. However, the content exceeding 0.0030% deteriorates the base material and the HAZ toughness. For this reason, when contained, B is preferably limited to a range of 0.0003% to 0.0030%. In addition, More preferably, it is 0.0006 to 0.0020%.

Ca:0.0005〜0.0050%、REM:0.0010〜0.0050%の1種または2種
Ca、REMはいずれも、硫化物の形態制御を介して母材の靭性向上および延性向上に寄与する。また、微細な硫化物粒子を鋼中に分散させた場合には、フェライト変態核として作用することによってHAZ靱性の向上にも寄与する。これらの効果を得るためには、Caでは少なくとも0.0005%、REMは少なくとも0.010%を含有することが好ましいが、Ca、REMをいずれも0.0050%を超えて含有すると、過剰な介在物が生成し、逆に靱性が低下する場合がある。このため、含有する場合には、Caは0.0005〜0.0050%、REMは0.0010〜0.0050%の範囲に限定することが好ましい。
Ca: 0.0005-0.0050%, REM: 0.0010-0.0050% of 1 type or 2 types of Ca, REM are both improved toughness and ductility of the base metal through the control of the form of sulfide. Contribute to. Further, when fine sulfide particles are dispersed in the steel, it acts as a ferrite transformation nucleus, thereby contributing to improvement of HAZ toughness. In order to obtain these effects, Ca preferably contains at least 0.0005% and REM contains at least 0.010%. However, if both Ca and REM contain more than 0.0050%, excessive amounts are required. Inclusions are produced, and the toughness may be reduced. For this reason, when it contains, it is preferable to limit Ca to 0.0005 to 0.0050% and REM to 0.0010 to 0.0050%.

なお、Ca、REMを含有する場合には、硫化物の形態制御作用を確保するために、下式で定義される炭素当量ACRが0.2〜0.8を満足するように、調整することが好ましい。   When Ca and REM are contained, the carbon equivalent ACR defined by the following formula should be adjusted to satisfy 0.2 to 0.8 in order to ensure the sulfide morphology control action. Is preferred.

ACRをこの範囲に調整すると、圧延時に鋼中に微細なCaおよび/またはREMの硫酸化物(オキシサルファイド)が分散し、さらに溶接後の冷却時に介在物(Caおよび/またはREMの硫酸化物)表面にMnSが析出する。   When the ACR is adjusted within this range, fine Ca and / or REM sulfates (oxysulfide) are dispersed in the steel during rolling, and the inclusions (Ca and / or REM sulfates) surface during cooling after welding. MnS precipitates on the surface.

このような複合介在物・析出物が粒内フェライトの生成サイトとして機能して、溶接ボンド部付近のミクロ組織が低靭性の上部ベイナイトで占められることを防止し、靭性を向上させる。   Such composite inclusions / precipitates function as intragranular ferrite formation sites, preventing the microstructure near the weld bond portion from being occupied by low toughness upper bainite and improving toughness.

ACRが0.2未満ではCa、REM量が不足し、所望の複合介在物・析出物を生成させることができないうえ、母材および溶接熱影響部靭性に有害なMnSが増加する。一方、ACRが0.8を超えると、ほとんどのSはCaおよび/またはREMの介在物中に取り込まれ、介在物表面に析出するMnSが不足し、介在物が粒内フェライト生成サイトとして十分に機能しなくなる。このため、ACRは0.2〜0.8の範囲に限定することが好ましい。   If the ACR is less than 0.2, the amount of Ca and REM is insufficient, so that desired composite inclusions / precipitates cannot be generated, and MnS harmful to the base material and weld heat affected zone toughness increases. On the other hand, when the ACR exceeds 0.8, most of S is taken into the inclusions of Ca and / or REM, MnS precipitated on the inclusion surface is insufficient, and the inclusions are sufficient as intragranular ferrite formation sites. Stops functioning. For this reason, it is preferable to limit ACR to the range of 0.2-0.8.

ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S)
(ここで、Ca、REM、O、Sは各元素の含有量(質量%))
なお、上記した成分以外の残部は、Feおよび不可避的不純物からなる。なお、不可避的不純物としては、O:0.005%以下が許容できる。
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S)
(Ca, REM, O, and S are the contents of each element (mass%))
The balance other than the above components is composed of Fe and inevitable impurities. As an inevitable impurity, O: 0.005% or less is acceptable.

[ミクロ組織]
本発明では、鋼板の少なくとも表層部(鋼板表面から板厚方向に1mm〜5mmの領域で、以降同じとする)のミクロ組織を平均結晶粒径が4.0〜18.0μmのフェライトと、硬質相としてパーライト、ベイナイトおよびマルテンサイトの1種または2種以上からなり、フェライト面積率が45%〜70%の複相組織とする。
[Microstructure]
In the present invention, the microstructure of at least the surface layer portion of the steel sheet (in the region from 1 mm to 5 mm in the thickness direction from the steel sheet surface, hereinafter the same), the ferrite having an average crystal grain size of 4.0 to 18.0 μm, and hard The phase is composed of one or more of pearlite, bainite and martensite, and has a multiphase structure with a ferrite area ratio of 45% to 70%.

降伏比75%以下の低降伏比と降伏強さ385MPa以上、引張強さ550MPa以上の高強度とを兼備させるために、鋼板の最表層(鋼板表面から板厚方向に1mm未満の領域)を除き、軟質相であるフェライトと、パーライト、ベイナイトおよびマルテンサイトの1種または2種以上からなる硬質相からなる複相組織とする。   Except for the outermost layer of steel plate (region less than 1mm from the steel plate surface to the plate thickness direction) in order to combine a low yield ratio of 75% or less with a yield strength of 385 MPa or higher and a tensile strength of 550 MPa or higher. A multiphase structure consisting of a soft phase of ferrite and a hard phase composed of one or more of pearlite, bainite and martensite.

軟質相であるフェライトと前記硬質相を組み合わせた複相組織とすることにより、優れた延性と所望の高強度、さらに低降伏比とを両立させることができる。フェライト面積率は、YR75%以下を達成するため45%以上とし、また降伏強さ385MPa以上、引張強さ550MPa以上を確保するため70%以下とする。   By forming a multiphase structure in which the soft phase ferrite and the hard phase are combined, it is possible to achieve both excellent ductility, desired high strength, and a low yield ratio. The ferrite area ratio is 45% or more in order to achieve YR of 75% or less, and 70% or less to ensure a yield strength of 385 MPa or more and a tensile strength of 550 MPa or more.

硬質相は、パーライト、ベイナイト、マルテンサイトのうちから選ばれた1種または2種以上とする。硬質相の種類とそれらの分率は、所望の強度と靭性、さらには化学成分や板厚によって適宜選択することができる。   The hard phase is one or more selected from pearlite, bainite, and martensite. The types of hard phases and their fractions can be appropriately selected depending on the desired strength and toughness, as well as chemical components and plate thickness.

フェライトの平均結晶粒径は、靭性および降伏比や伸びに大きく影響する。フェライト粒径が4.0μm未満では、降伏比が急激に上昇し、均一伸びが低下し、塑性変形能が大きく低下する。一方、フェライト粒径が18.0μmを超えて粗大化すると、靱性が低下し、脆性破壊が発生しやすくなるため、4.0〜18.0μmの範囲に限定した。なお、好ましくは7.0〜14.0μmである。平均結晶粒径の求め方は実施例において説明する。   The average crystal grain size of ferrite greatly affects toughness, yield ratio and elongation. When the ferrite particle size is less than 4.0 μm, the yield ratio increases rapidly, the uniform elongation decreases, and the plastic deformability greatly decreases. On the other hand, when the ferrite grain size exceeds 18.0 μm and becomes coarse, the toughness is lowered and brittle fracture is likely to occur. Therefore, the ferrite grain size is limited to the range of 4.0 to 18.0 μm. In addition, Preferably it is 7.0-14.0 micrometers. The method for obtaining the average crystal grain size will be described in Examples.

鋼板の表層部が上述したミクロ組織の場合、鋼板表層部の塑性変形能が大きく向上するので、冷間曲げ加工を施した後に、鋼板表層部の靭性・延性の低下が抑制され、地震による変形時に溶接止端部などからの延性亀裂の発生が抑制される。   When the surface layer portion of the steel sheet has the above-described microstructure, the plastic deformability of the steel sheet surface layer portion is greatly improved, so that after the cold bending process, the deterioration of the toughness and ductility of the steel sheet surface layer portion is suppressed, and deformation due to earthquake Occurrence of ductile cracks from the weld toe is sometimes suppressed.

[板厚方向の硬さ分布]
鋼板の表層部の平均硬さが225HV以下で、該表層部と板厚中央部(板厚中央位置を中心に±2mmの領域で、以降同じとする)との硬度差が60HV以下である板厚方向の硬さ分布とする。
[Hardness distribution in the thickness direction]
The average hardness of the surface layer portion of the steel sheet is 225 HV or less, and the hardness difference between the surface layer portion and the plate thickness central portion (the region of ± 2 mm centered on the plate thickness center position is the same hereinafter) is 60 HV or less. The hardness distribution in the thickness direction.

鋼板表層部の硬さが225HVを超えると、冷間曲げ加工を施した後に、鋼板表層部の硬さがさらに増加し、鋼板表層部の靭性・延性が著しく低下し、建築構造物の柱−梁接合部などの部材(T継手、十字継手)で、地震による変形時に溶接止端部など表層部から亀裂を発生しやすくなる。   If the hardness of the steel sheet surface layer part exceeds 225 HV, after the cold bending process, the hardness of the steel sheet surface layer part further increases, the toughness and ductility of the steel sheet surface layer part significantly decrease, and the column of the building structure In members such as beam joints (T joints and cross joints), cracks are likely to occur from the surface layer parts such as weld toes at the time of deformation due to an earthquake.

また、板厚全体で所定の強度を確保するため、板厚方向の硬度差はできるだけ少ないことが望ましい。表層部と板厚中央部との硬度差が60HV超えて大きくなると、地震等による変形時に、溶接接合部の形状によっては、相対的に強度の低い板厚中央部に変形が集中し、早期に破壊が発生する場合がある。このような破壊を防止して必要な部材性能を確保するために、表層部と板厚中央部の硬度差を60HV以下、なお、好ましくは硬度差55HV以下、より好ましくは50HV以下とする。板厚中央部の硬さは175〜200HVとすることが所望の高強度を確保するために好ましい。硬さ試験方法は実施例において説明する。   Further, in order to ensure a predetermined strength throughout the thickness, it is desirable that the difference in hardness in the thickness direction is as small as possible. When the hardness difference between the surface layer and the plate thickness center exceeds 60 HV, the deformation concentrates on the plate thickness center with relatively low strength depending on the shape of the welded joint during deformation due to an earthquake, etc. Destruction may occur. In order to prevent such destruction and ensure necessary member performance, the hardness difference between the surface layer portion and the plate thickness center portion is 60 HV or less, preferably, the hardness difference is 55 HV or less, more preferably 50 HV or less. The hardness at the center of the plate thickness is preferably 175 to 200 HV in order to ensure the desired high strength. The hardness test method will be described in Examples.

本発明に係る厚鋼板は、上記組成の鋼素材に、熱間圧延を施し厚鋼板とする圧延工程と、該圧延工程に引続き、一段冷却と一段冷却を停止し、復熱後に行う二段冷却とからなる二段階の加速冷却を行う冷却工程とを備えた製造方法で製造する。   The steel sheet according to the present invention is a steel sheet having the above composition, which is subjected to hot rolling to form a steel sheet, and subsequently to the rolling process, one-stage cooling and one-stage cooling are stopped, and two-stage cooling performed after reheating. And a cooling process for performing two-stage accelerated cooling consisting of:

鋼素材の製造方法は、特に限定する必要はなく、常用の溶製方法、鋳造方法がいずれも適用できるが、上記した組成の溶鋼を、転炉、電気炉、真空溶解炉等で溶製し、脱酸処理や脱ガスプロセスを経て、連続鋳造法などで鋼素材(スラブ)とすることが好ましい。   The manufacturing method of the steel material is not particularly limited, and any conventional melting method and casting method can be applied, but the molten steel having the above composition is melted in a converter, electric furnace, vacuum melting furnace or the like. The steel material (slab) is preferably obtained by a continuous casting method through a deoxidation treatment or a degassing process.

得られた鋼素材(スラブ)は、まず加熱され、熱間圧延されて厚鋼板となる圧延工程を施される。圧延工程では、鋼素材を加熱温度:1050〜1200℃に加熱し、表面温度で950℃以下の温度域での累積圧下量が30%以上で、圧延終了温度が表面温度で900℃以下Ar3変態点以上とする熱間圧延を施し、所定板厚の厚鋼板とする。   The obtained steel material (slab) is first heated and subjected to a rolling process that is hot-rolled into a thick steel plate. In the rolling process, the steel material is heated to a heating temperature of 1050 to 1200 ° C., the cumulative reduction amount in the temperature range of 950 ° C. or less at the surface temperature is 30% or more, and the rolling end temperature is 900 ° C. or less at the surface temperature. Hot rolling at a point or more is performed to obtain a thick steel plate having a predetermined thickness.

加熱温度:1050〜1200℃
加熱温度が1050℃未満では、強度が低下しやすく、一方、1200℃を超えると、組織が粗大化して得られる靱性が低下したり、焼入性が増加しすぎて、表層硬さが増加しやすくなる。このため、鋼素材の加熱温度は1050℃〜1200℃の範囲に限定した。なお、好ましくは1080℃〜1150℃である。
Heating temperature: 1050-1200 ° C
If the heating temperature is less than 1050 ° C., the strength tends to decrease. On the other hand, if it exceeds 1200 ° C., the toughness obtained by coarsening the structure decreases, the hardenability increases excessively, and the surface hardness increases. It becomes easy. For this reason, the heating temperature of the steel material was limited to the range of 1050 ° C to 1200 ° C. In addition, Preferably it is 1080 to 1150 degreeC.

表面温度で950℃以下の温度域での累積圧下量:30%以上
ミクロ組織を適度に微細化するため、鋼板の表面温度が950℃以下の温度域で制御圧延を行う。該温度域での累積圧下量が30%未満では、組織が粗大化し、また焼入性が増加しすぎて、所望の靭性、表層硬さを確保できなくなる。このため、表面温度で950℃以下の温度域での累積圧下量を30%以上に限定した。なお、好ましくは35%以上である。
Cumulative reduction in surface temperature range of 950 ° C. or lower: 30% or more Control rolling is performed in a temperature range where the surface temperature of the steel sheet is 950 ° C. or lower in order to refine the microstructure appropriately. When the cumulative reduction amount in the temperature range is less than 30%, the structure becomes coarse and hardenability increases too much, so that desired toughness and surface hardness cannot be ensured. For this reason, the cumulative reduction amount in the temperature range of 950 ° C. or less at the surface temperature is limited to 30% or more. In addition, Preferably it is 35% or more.

圧延終了温度:表面温度で900℃以下Ar3変態点以上
圧延終了温度が表面温度で900℃を超えると、組織が粗大化し、焼入性が増加しすぎて、所望の靭性、表層硬さを確保できなくなる。一方、圧延終了温度が表面温度でAr3変態点未満では、圧延中あるいは圧延直後にフェライトが生成し、粗大化して、表層部の靱性が低下する。このため、圧延終了温度は表面温度で900℃以下Ar3温度以上に限定した。なお、好ましくは880〜780℃である。
Rolling end temperature: 900 ° C. or less at the surface temperature and Ar3 transformation point or higher If the rolling end temperature exceeds 900 ° C. at the surface temperature, the structure becomes coarse and hardenability increases, ensuring the desired toughness and surface hardness. become unable. On the other hand, if the rolling end temperature is less than the Ar3 transformation point at the surface temperature, ferrite is generated during rolling or immediately after rolling and becomes coarse, and the toughness of the surface layer portion decreases. For this reason, the rolling end temperature was limited to 900 ° C. or lower and Ar 3 temperature or higher as the surface temperature. In addition, Preferably it is 880-780 degreeC.

なお、Ar3変態点は、下記式を用いて算出した値を用いるものとする。   As the Ar3 transformation point, a value calculated using the following formula is used.

Ar3変態点(℃)=900−332C+6Si−77Mn−20Cu−50Ni−18Cr−68Mo
(式において、C、Si、Mn、Cu、Ni、Cr、Mo:各元素の含有量(質量%)で、上記式で記載された元素が含有されない場合には、当該元素を0として計算するものとする。)
圧延後、加速冷却する冷却工程を施す。冷却工程は、第一段冷却と、冷却を停止し復熱させる過程と、第二段冷却とからなる。第一段冷却で、表層部を過冷却したのち復熱させ、第二段冷却の開始までの時間(冷却停止時間)を利用して、表層部のフェライト変態を進行させて所望の表層ミクロ組織を得る。第二段冷却は、第一段冷却後に、未変態である部分をパーライト、ベイナイト、マルテンサイト等の硬質相とするために行う。未変態部分を硬質相とすることにより、最終組織を(フェライト+(パーライト,ベイナイトおよび/またはマルテンサイト))とすることができ、所望の高強度、低降伏比を実現できる。
Ar3 transformation point (° C) = 900-332C + 6Si-77Mn-20Cu-50Ni-18Cr-68Mo
(In the formula, C, Si, Mn, Cu, Ni, Cr, Mo: The content (mass%) of each element, and when the element described in the above formula is not contained, the element is calculated as 0 Suppose)
After rolling, a cooling step for accelerated cooling is performed. The cooling step includes first-stage cooling, a process of stopping cooling and returning to heat, and second-stage cooling. In the first stage cooling, the surface layer part is supercooled and then reheated, and the time until the start of the second stage cooling (cooling stop time) is used to advance the ferrite transformation of the surface layer part to achieve the desired surface layer microstructure. Get. The second stage cooling is carried out after the first stage cooling so that the untransformed portion becomes a hard phase such as pearlite, bainite, martensite or the like. By making the untransformed portion a hard phase, the final structure can be (ferrite + (pearlite, bainite and / or martensite)), and a desired high strength and low yield ratio can be realized.

[第一段冷却]
第一段冷却は、表面温度でAr3変態点以上の温度から冷却を開始し、板厚(t)の1/4位置の平均冷却速度で2℃/s以上の冷却速度で加速冷却し、表面温度が(Ar3変態点−100℃)以下550℃以上、好ましくはかつ表面と板厚中央位置との温度差が150℃以上、となる時点で、加速冷却を停止する。
[First stage cooling]
In the first stage cooling, cooling is started from a temperature equal to or higher than the Ar3 transformation point at the surface temperature, and accelerated cooling is performed at a cooling rate of 2 ° C./s or higher at an average cooling rate at 1/4 position of the plate thickness (t). Accelerated cooling is stopped when the temperature reaches (Ar3 transformation point−100 ° C.) or less, 550 ° C. or more, and preferably the temperature difference between the surface and the plate thickness center position is 150 ° C. or more.

第一段冷却の開始温度:表面温度でAr3変態点以上
第一段冷却の開始温度が、Ar3変態点未満では、加速冷却開始前にフェライトが生成し、粗大化するため、表層部のフェライト粒の微細化が達成できなくなる。このため、第一段冷却の開始温度をAr3変態点以上に限定した。
First stage cooling start temperature: Ar3 transformation point or higher at the surface temperature If the first stage cooling start temperature is less than the Ar3 transformation point, ferrite is generated and coarsened before the start of accelerated cooling. Cannot be achieved. For this reason, the start temperature of the first stage cooling is limited to the Ar3 transformation point or higher.

第一段冷却の冷却速度:板厚(t)の1/4位置の平均冷却速度で2℃/s以上
冷却速度が2℃/s未満では、冷却が遅く、冷却中に粗く靭性の低いフェライト粒が生成する場合がある。このため、第一段冷却の冷却速度を、板厚(t)の1/4位置の平均冷却速度で2℃/s以上に限定した。なお、第一段冷却の冷却速度の上限はとくに限定する必要はなく、板厚、冷却装置の能力によってほぼ決定され、板厚:60mmでは概ね5℃/s程度以上となる。「板厚(t)の1/4位置の平均冷却速度」とは、板厚(t)の1/4位置における加速冷却開始から終了までの平均の冷却速度をいう。
Cooling rate of the first stage cooling: 2 ° C./s or more at an average cooling rate of 1/4 position of the plate thickness (t) When the cooling rate is less than 2 ° C./s, the cooling is slow and the ferrite is coarse and low toughness during cooling. Grain may form. For this reason, the cooling rate of the first stage cooling was limited to 2 ° C./s or more as the average cooling rate at the 1/4 position of the plate thickness (t). The upper limit of the cooling rate of the first stage cooling is not particularly limited, and is almost determined by the plate thickness and the capacity of the cooling device, and is about 5 ° C./s or more when the plate thickness is 60 mm. The “average cooling rate at the 1/4 position of the plate thickness (t)” refers to the average cooling rate from the start to the end of the accelerated cooling at the 1/4 position of the plate thickness (t).

第一段冷却の冷却停止温度:表面温度で(Ar3変態点−100℃)以下550℃以上
本発明における第一段冷却では、表層部とそれより内部との温度差が大きくなるように冷却し、第一段冷却停止後の復熱と、復熱後、第二段冷却を開始するまでの時間(以降、保持時間と呼ぶ)に、表層部にフェライトを生成させる。
Cooling stop temperature of the first stage cooling: (Ar3 transformation point −100 ° C.) or less at the surface temperature 550 ° C. or more In the first stage cooling in the present invention, the cooling is performed so that the temperature difference between the surface layer portion and the inside becomes larger. Then, ferrite is generated in the surface layer portion during the recuperation after stopping the first stage cooling and the time after the recuperation until starting the second stage cooling (hereinafter referred to as holding time).

冷却停止温度が、表面温度で(Ar3変態点−100℃)を超えると、その後の復熱温度が高すぎて、表層部におけるフェライト生成が不十分となる。一方、冷却停止温度が550℃未満では、表層部の温度が低温となりすぎて、復熱後の冷却中に相変態がほぼ完了し、表層部がベイナイトやマルテンサイトなどの硬質相主体の組織となる。このため、第一段冷却の冷却停止温度は表面温度で(Ar3変態点−100℃)以下550℃以上とする。第一段冷却の冷却停止温度は、好ましくは650〜550℃である。   When the cooling stop temperature exceeds the surface temperature (Ar3 transformation point −100 ° C.), the subsequent recuperation temperature is too high, and ferrite formation in the surface layer becomes insufficient. On the other hand, if the cooling stop temperature is less than 550 ° C., the temperature of the surface layer portion becomes too low, and phase transformation is almost completed during cooling after recuperation, and the surface layer portion is composed of a structure mainly composed of a hard phase such as bainite and martensite. Become. For this reason, the cooling stop temperature of the first stage cooling is set to the surface temperature (Ar3 transformation point−100 ° C.) or lower and 550 ° C. or higher. The cooling stop temperature of the first stage cooling is preferably 650 to 550 ° C.

[復熱および第二段冷却]
第一段冷却を停止したのち、フェライトを生成させるために復熱後、第二段冷却を開始する。復熱は、復熱後に鋼板表面温度が(Ar3変態点−20℃)以下600℃以上となり、且つ所定の保持時間を満足するまで行う。
[Recuperation and second stage cooling]
After stopping the first stage cooling, the second stage cooling is started after reheating to generate ferrite. The recuperation is performed until the steel sheet surface temperature is (Ar3 transformation point−20 ° C.) or lower and 600 ° C. or higher after the reheating and a predetermined holding time is satisfied.

鋼板表面温度が600℃未満では、表層部において強度および降伏比が比較的高い針状フェライトやベイナイトが生成するので、表層部の伸びの低下や降伏比の上昇などが生じ、変形能の低下を招く。一方、表面温度で(Ar3変態点−20℃)を超えると、復熱後に相変態が進行せず、表層部におけるフェライト生成が不十分となるため、(Ar3変態点−20℃)以下600℃以上となるまで保持する。   If the surface temperature of the steel sheet is less than 600 ° C., acicular ferrite and bainite having a relatively high strength and yield ratio are formed in the surface layer portion, resulting in a decrease in elongation of the surface layer portion and an increase in yield ratio, resulting in a decrease in deformability. Invite. On the other hand, if the surface temperature exceeds (Ar3 transformation point −20 ° C.), phase transformation does not proceed after recuperation, and ferrite formation in the surface layer portion becomes insufficient, so (Ar3 transformation point −20 ° C.) or lower 600 ° C. Hold until above.

保持時間(復熱後、鋼板の表面温度が極大値をとった時点から、第二段冷却の開始までの時間)は下記(3)式を満たす時間t1(秒)以上、下記(4)式を満たす時間t2(秒)以内とする。   The holding time (the time from when the surface temperature of the steel sheet takes the maximum value after recuperation until the start of the second stage cooling) is equal to or longer than the time t1 (seconds) that satisfies the following equation (3), and the following equation (4): Within a time t2 (seconds) that satisfies

保持時間がt1(秒)未満であると、鋼板表層部に析出するフェライトが50%未満になり、鋼板の表層軟化、およびYR≦80%を達成できなくなる。また、保持時間がt2(秒)を超えると、鋼板の表層部に析出するフェライトが70%を超えてしまい、鋼板が軟化し、TS強度が550MPa未満となるため保持時間はt1(秒)以上t2(秒)以内とする。   When the holding time is less than t1 (seconds), the ferrite precipitated on the surface layer portion of the steel sheet is less than 50%, and the surface layer softening of the steel sheet and YR ≦ 80% cannot be achieved. If the holding time exceeds t2 (seconds), the ferrite precipitated on the surface layer of the steel sheet exceeds 70%, the steel sheet softens, and the TS strength becomes less than 550 MPa, so the holding time is t1 (seconds) or more. Within t2 (seconds).

このように、復熱を行うと、第二段冷却開始時点で表面と板厚中央の温度差が60℃以下となり、第二段冷却終了時点で、表層部のフェライト面積分率として45〜70%が得られ、且つ表面と板厚中央の板厚方向のフェライト生成量の差を小さくし、表層部と板厚中央部との硬度差を60HV以下とすることができる。   Thus, when reheating is performed, the temperature difference between the surface and the center of the plate thickness is 60 ° C. or less at the start of the second stage cooling, and the ferrite area fraction of the surface layer portion is 45 to 70 at the end of the second stage cooling. %, The difference in the amount of ferrite produced in the thickness direction between the surface and the center of the plate thickness can be reduced, and the hardness difference between the surface layer portion and the plate thickness center portion can be reduced to 60 HV or less.

第二段冷却は、板厚(t)の1/4位置で2℃/s以上の平均冷却速度で、該第二段冷却を停止した後の復熱で表面温度が600℃以下400℃以上になる冷却停止温度まで加速冷却する。   The second stage cooling is an average cooling rate of 2 ° C./s or more at the 1/4 position of the plate thickness (t), and the surface temperature is 600 ° C. or less and 400 ° C. or more by recuperation after stopping the second stage cooling. Accelerate cooling to the cooling stop temperature.

第二段冷却の冷却速度:板厚(t)の1/4位置の平均冷却速度で2℃/s以上
未変態部分を硬質相とするために、第二段冷却では、2℃/s以上、好ましくは8℃/s以上で冷却する。冷却速度が2℃/s未満では、硬質相への変態量が低下し、所望の高強度、低降伏比を実現できなくなる。
Cooling rate of the second stage cooling: 2 ° C./s or more at an average cooling rate at 1/4 position of the thickness (t) In order to make the untransformed part a hard phase, 2 ° C./s or more in the second stage cooling The cooling is preferably performed at 8 ° C./s or more. If the cooling rate is less than 2 ° C./s, the amount of transformation to the hard phase decreases, and the desired high strength and low yield ratio cannot be realized.

第二段冷却の冷却停止温度:冷却を停止した後の復熱で表面温度が600℃以下400℃以上になる温度
第二段冷却の冷却停止温度が、第二段冷却の冷却停止後の復熱で表面温度が600℃超えとなる温度では、硬質相への変態量が低下したり、自己焼戻しによって強度が低下し、所望の高強度を確保できなくなる。一方、復熱で表面温度が400℃未満となるような冷却停止温度では、硬質相硬さが高くなりすぎて靱性が低下する。このため、第二段冷却の冷却停止温度は、冷却を停止した後の復熱で表面温度が600℃以下400℃以上になる温度に限定した。復熱後の温度は、加速冷却停止時の板厚(t)の1/2位置の温度に依存するので各種伝熱計算から予測することができる。
Cooling stop temperature of the second stage cooling: Temperature at which the surface temperature becomes 600 ° C. or lower and 400 ° C. or higher by the recuperation after stopping the cooling. The cooling stop temperature of the second stage cooling is the recovery temperature after the cooling stop of the second stage cooling. At a temperature at which the surface temperature exceeds 600 ° C. due to heat, the amount of transformation to the hard phase decreases, or the strength decreases due to self-tempering, making it impossible to ensure the desired high strength. On the other hand, at the cooling stop temperature at which the surface temperature is less than 400 ° C. due to recuperation, the hard phase hardness becomes too high and the toughness is lowered. For this reason, the cooling stop temperature of the second stage cooling is limited to a temperature at which the surface temperature becomes 600 ° C. or lower and 400 ° C. or higher by reheating after cooling is stopped. Since the temperature after recuperation depends on the temperature at the half position of the plate thickness (t) when the accelerated cooling is stopped, it can be predicted from various heat transfer calculations.

第二段冷却後、強度および靭性の調整を目的として、焼戻工程を施してもよい。焼戻しは、400℃以上700℃以下の温度で行うことが好ましい。焼戻温度が400℃未満では、所望の効果を期待できない。一方、700℃を超える温度では、強度低下が著しくなる。   After the second stage cooling, a tempering step may be performed for the purpose of adjusting strength and toughness. Tempering is preferably performed at a temperature of 400 ° C. or higher and 700 ° C. or lower. If the tempering temperature is less than 400 ° C., the desired effect cannot be expected. On the other hand, when the temperature exceeds 700 ° C., the strength is significantly reduced.

本発明では、第一段冷却と第二段冷却の両方、またはいずれかのみを、1回の加速冷却からなる冷却に代えて、冷却停止とその後の復熱とを挟んで、複数回繰り返す冷却としてもよい。   In the present invention, both the first-stage cooling and the second-stage cooling, or only one of them, are replaced with cooling consisting of one accelerated cooling, and cooling is repeated a plurality of times with a cooling stop and subsequent recuperation. It is good.

加速冷却を複数回に分割することにより、表層と内部との温度差を、過度に大きくすることなく、目的の温度まで冷却することが可能となる。また、複数回繰り返す中で所期の冷却効果を得ればよいことから、冷却温度制御の選択肢が拡大でき、冷却温度制御の精度を向上させることができる。図7にこのような冷却を行った場合の鋼板温度の履歴の一例を模式的に示す。なお図において保持時間は上述した復熱の保持時間を示す。   By dividing the accelerated cooling into a plurality of times, it is possible to cool to the target temperature without excessively increasing the temperature difference between the surface layer and the inside. In addition, since it is sufficient to obtain a desired cooling effect while it is repeated a plurality of times, options for cooling temperature control can be expanded, and the accuracy of cooling temperature control can be improved. FIG. 7 schematically shows an example of the history of the steel sheet temperature when such cooling is performed. In the drawing, the holding time indicates the holding time of the above-described recuperation.

第一段冷却を複数回の冷却によるものとする場合、最初の加速冷却(第1回冷却)の開始温度は一回の場合と同様の理由により表面温度でAr3変態点以上とする。   When the first stage cooling is performed by a plurality of times of cooling, the starting temperature of the first accelerated cooling (first cooling) is set to the Ar3 transformation point or more at the surface temperature for the same reason as in the case of one time.

加速冷却の冷却速度は第一段冷却における最初の冷却と最後の冷却までの板厚(t)の1/4位置の平均冷却速度とし、一回の場合と同様の理由により2℃/s以上とする。   The cooling rate of accelerated cooling is the average cooling rate at 1/4 position of the plate thickness (t) from the first cooling to the last cooling in the first stage cooling, and it is 2 ° C./s or more for the same reason as in the case of one time. And

板厚(t)の1/4位置の平均冷却速度は、板厚(t)の1/4位置における加速冷却開始から終了までの平均の冷却速度で図7のA点からB点までの平均の冷却速度とする。A点は、板厚1/4t位置における温度が表面の冷却開始温度に等しくなった時点であり、B点は、第一段冷却における最後の加速冷却を停止した時点である。   The average cooling rate at the 1/4 position of the plate thickness (t) is the average cooling rate from the start to the end of accelerated cooling at the 1/4 position of the plate thickness (t), and the average from point A to point B in FIG. Cooling rate. Point A is the time when the temperature at the plate thickness 1 / 4t position becomes equal to the surface cooling start temperature, and point B is the time when the last accelerated cooling in the first stage cooling is stopped.

第一段冷却における複数回の加速冷却において、冷却停止温度が表層部で550℃未満となると、冷却中に、ベイナイト、マルテンサイト変態が生じて、表層部が硬質化するため、すべての加速冷却の冷却停止温度を550℃以上とする。   In the multiple times of accelerated cooling in the first stage cooling, when the cooling stop temperature is less than 550 ° C. in the surface layer part, bainite and martensite transformation occurs during cooling, and the surface layer part becomes hard, so all accelerated cooling The cooling stop temperature is set to 550 ° C. or higher.

第一段冷却は、表層部と内部との温度差がある程度生じるように冷却し、冷却停止後の復熱、及び保持時間の間に、表層部にフェライトを生成させることを目的とするので、複数回の加速冷却の全てが、冷却停止温度が表面温度で(Ar3変態点−100℃)を超えると、その後の復熱時に、鋼板温度が高くなりすぎて、表層部でのフェライト生成が不十分となる。このため、複数回の加速冷却のうち、少なくとも1回の冷却停止温度を(Ar3変態点−100℃)以下とする。   The first stage cooling is performed so that a temperature difference between the surface layer part and the inside is generated to some extent, and the purpose is to generate ferrite in the surface layer part during the recuperation after the cooling stop and the holding time. In all of the multiple times of accelerated cooling, if the cooling stop temperature exceeds the surface temperature (Ar3 transformation point −100 ° C.), the steel sheet temperature becomes too high at the time of subsequent reheating, and ferrite formation at the surface layer portion is not possible. It will be enough. For this reason, at least one cooling stop temperature is set to (Ar3 transformation point−100 ° C.) or lower among a plurality of times of accelerated cooling.

第二段冷却の冷却を、冷却停止とその後の復熱とを挟んで、複数回繰り返す冷却とする場合、最初の加速冷却(第1回冷却)は一回の場合と同様に復熱後の温度と時間を満足させた後に開始する。   When the cooling of the second stage cooling is repeated a plurality of times with the cooling stop and the subsequent recuperation interposed therebetween, the first accelerated cooling (first cooling) is performed after the recuperation as in the case of one time. Start after satisfying temperature and time.

第二段冷却を構成する複数回の加速冷却の冷却速度は第一段冷却における最初の冷却と最後の冷却までの板厚1/4t位置の平均冷却速度とし、一回の場合と同様の理由により2℃/s以上とする。   The cooling rate of the multiple times of the accelerated cooling that constitutes the second stage cooling is the average cooling rate at the position of 1 / 4t thickness until the first cooling and the last cooling in the first stage cooling, and the same reason as the one-time cooling To 2 ° C./s or more.

板厚(t)の1/4位置の平均冷却速度は、板厚(t)の1/4位置における加速冷却開始から終了までの平均の冷却速度で図7のC点からD点までの平均の冷却速度とする。C点は、板厚(t)の1/4位置における温度が表面の冷却開始温度に等しくなった時点であり、D点は、第二段冷却における最後の加速冷却を停止した時点である。   The average cooling rate at the 1/4 position of the plate thickness (t) is the average cooling rate from the start to the end of the accelerated cooling at the 1/4 position of the plate thickness (t), and the average from point C to point D in FIG. Cooling rate. Point C is the time when the temperature at the 1/4 position of the plate thickness (t) becomes equal to the surface cooling start temperature, and point D is the time when the last accelerated cooling in the second stage cooling is stopped.

第二段冷却の最終冷却は、一回の場合と同様の理由により、冷却停止後の復熱で表面温度が600℃以下400℃以上になるような冷却停止温度まで冷却する。   The final cooling of the second stage cooling is performed for the same reason as in the case of one time, and is cooled to a cooling stop temperature at which the surface temperature becomes 600 ° C. or lower and 400 ° C. or higher by recuperation after cooling stop.

所望する冷却温度制御の精度に応じて、第一段冷却と第二段冷却の両方、またはいずれかのみを複数回繰り返す冷却とする。以下、実施例を用いて更に本発明を詳細に説明する。   Depending on the accuracy of desired cooling temperature control, both the first-stage cooling and the second-stage cooling, or only one of them, is repeated multiple times. Hereinafter, the present invention will be described in more detail with reference to examples.

表1に示す組成を有する鋼素材に、表2に示す圧延工程、冷却工程を施し、板厚:40mmの厚鋼板とした。冷却工程では、第一段冷却の冷却停止−復熱後、第二段冷却を施した。各工程における、鋼板温度は、赤外線放射温度計で表面温度を測定し、板厚1/4t位置の温度、板厚中央温度を種々の伝熱計算法を用いて算出した。   The steel material having the composition shown in Table 1 was subjected to the rolling step and the cooling step shown in Table 2 to obtain a thick steel plate having a plate thickness of 40 mm. In the cooling step, the second stage cooling was performed after the cooling stop of the first stage cooling-recuperation. The steel plate temperature in each step was measured by measuring the surface temperature with an infrared radiation thermometer, and calculating the temperature at the plate thickness ¼ t position and the plate thickness central temperature using various heat transfer calculation methods.

得られた厚鋼板について、組織観察、硬さ試験、引張試験、衝撃試験を実施した。試験方法は次の通りとした。
(1)組織観察
板厚全厚の組織観察用試験片のL方向断面を研磨、ナイタール腐食後、表層部と板厚中央部を光学顕微鏡(倍率:400倍)または走査型電子顕微鏡(倍率:2000倍)を用いて、ミクロ組織を各3視野以上観察し、撮像して画像解析により、組織の種類、およびフェライトの組織分率(面積率%)を求めた。
The obtained thick steel plate was subjected to structure observation, hardness test, tensile test, and impact test. The test method was as follows.
(1) Microstructure observation The cross section in the L direction of the specimen for structural observation of the full thickness of the plate is polished, and after nital corrosion, the surface layer portion and the central portion of the plate thickness are optical microscope (magnification: 400 times) or scanning electron microscope (magnification: 2000 times), the microstructure was observed for three or more visual fields, imaged, and image analysis was performed to determine the type of structure and the structure fraction (area ratio%) of ferrite.

また、表層部については、フェライトの平均結晶粒径(公称粒径という場合がある)を求めた。フェライトの公称粒径は、結晶粒の平均面積を求め、得られた結晶粒の平均面積の平方根とした。
(2)硬さ試験
板厚全厚の硬さ測定用試験片を採取し、ビッカース硬さ計を用いて、JIS Z 2244の規定に準拠して、板厚方向断面について、硬さ測定を行った。測定位置は、表層部、および板厚中央部とし、各領域で板厚方向に1mmピッチで、4点以上測定した。試験荷重(試験力)は1kg(9.8kN)とした。得られた硬さHVを算術平均し、その領域での平均硬さHVとした。
(3)引張試験
引張方向がL方向となるように、JIS Z 2201の規定に準拠して、JIS5号全厚引張試験片を採取し、JIS Z 2241の規定に準拠して、引張試験を実施し、引張特性(降伏強さYS、引張強さTS)を求めた。また、得られた測定値から、降伏比YR(=YS/TS×100%)を算出した。
(4)衝撃試験
板厚(t)の1/4位置および表面下1mm(試験片中央位置が表面下6mm)位置から、JIS Z 2242に準拠して、Vノッチ衝撃試験片を採取し、シャルピー衝撃試験を実施し、破面遷移温度vTrs(℃)を求めた。なお、vTrsが、−40℃以下である場合を靭性に優れるとした。
For the surface layer portion, the average crystal grain size of ferrite (sometimes referred to as a nominal grain size) was determined. The nominal grain size of the ferrite was determined as the square root of the average area of the crystal grains obtained by calculating the average area of the crystal grains.
(2) Hardness test Collect a test piece for hardness measurement of the full thickness of the plate, and measure the hardness of the cross section in the plate thickness direction using a Vickers hardness tester in accordance with the provisions of JIS Z 2244. It was. The measurement position was a surface layer portion and a plate thickness central portion, and four or more points were measured at a pitch of 1 mm in the plate thickness direction in each region. The test load (test force) was 1 kg (9.8 kN). The obtained hardness HV was arithmetically averaged to obtain the average hardness HV in that region.
(3) Tensile test JIS No. 5 full-thickness tensile test specimens are collected in accordance with the provisions of JIS Z 2201 so that the tensile direction is the L direction, and the tensile tests are performed in accordance with the provisions of JIS Z 2241. The tensile properties (yield strength YS, tensile strength TS) were determined. Moreover, the yield ratio YR (= YS / TS × 100%) was calculated from the obtained measured values.
(4) Impact test V-notch impact test specimens were collected from ¼ position of the plate thickness (t) and 1 mm below the surface (test specimen central position was 6 mm below the surface) according to JIS Z 2242, and Charpy. An impact test was performed to determine the fracture surface transition temperature vTrs (° C.). In addition, the case where vTrs is -40 degrees C or less was considered to be excellent in toughness.

(5)コラム曲げ試験
得られた厚鋼板を用いて、冷間プレス加工により、角形鋼管(プレスコラム)を作製した。角形鋼管(プレスコラム)の断面寸法は、500×500(mm)、長さは3250(mm)とし、シーム(継目)溶接は両面各1層のサブマージアーク溶接とした。
(5) Column bending test Using the obtained thick steel plate, a square steel pipe (press column) was produced by cold pressing. The cross-sectional dimensions of the square steel pipe (press column) were 500 × 500 (mm), the length was 3250 (mm), and the seam welding was submerged arc welding with one layer on each side.

図3にコラム曲げ試験の試験体6を示す。4面BOX柱3aの左右にSN490鋼板製通しダイアフラム(板厚40mm)2a、2aを介して角形鋼管(プレスコラム)1a、1aと取り付けた。各部材間の溶接は炭酸ガス溶接とした。   FIG. 3 shows a specimen 6 for the column bending test. Square steel pipes (press columns) 1a and 1a were attached to the left and right sides of the four-sided BOX column 3a through SN490 steel plate through diaphragms (plate thickness 40 mm) 2a and 2a. Welding between each member was carbon dioxide welding.

4面BOX柱の強度と剛性をプレスコラムに比べて十分高くすることにより、試験中にプレスコラム以外で塑性変形が生じないようにした。図4に試験体6におけるプレスコラム1a、ダイアフラム2a、4面BOX柱3aの溶接部近傍を拡大して示す。図において4a、5aは溶接部を示す。図示した試験体6を用いて、以下の要領でコラム曲げ試験を行った。   By making the strength and rigidity of the four-sided BOX column sufficiently higher than that of the press column, plastic deformation was prevented from occurring outside the press column during the test. FIG. 4 shows an enlarged view of the vicinity of the welded portion of the press column 1a, the diaphragm 2a, and the four-sided BOX column 3a in the test body 6. In the figure, reference numerals 4a and 5a denote welds. A column bending test was performed using the illustrated specimen 6 in the following manner.

試験体6の両端部を支持し、図5に示すように、試験体6の中央部に上下方向に正負の荷重を繰り返し負荷する、3点繰り返し曲げ試験(コラム曲げ試験)を実施した。荷重Pと変形量(回転角)θを測定し、図6に示すような荷重(モーメント、M)−変形量(回転角、θ)ヒステリシス曲線を作成した。   As shown in FIG. 5, a three-point repeated bending test (column bending test) was performed in which positive and negative loads were repeatedly applied in the vertical direction to the center of the test body 6 as shown in FIG. The load P and the deformation amount (rotation angle) θ were measured, and a load (moment, M) -deformation amount (rotation angle, θ) hysteresis curve as shown in FIG. 6 was created.

局部座屈または脆性破壊によって荷重(モーメント)が最大値から5%低下した時点を試験体の破壊とみなし、それまでの試験体の塑性回転角の合計(累積塑性回転角Σθpl)を求め、試験体の塑性変形能の指標として累積塑性変形倍率ηを求め、30以上である場合、構造部材の耐震性(塑性変形性能)に優れるとした。なお、ηは下式より算出される。
η=Σθpl/θp
但し、θp=(Pp/2)L/(3・E・I)+Pp/2/(G・Aw)
ここで、Pp:全塑性時荷重(N)=Mp/L、
L:コラムの片持ち長さ(ダイアフラムからコラム端支持点までの距離、3250mm)
E:ヤング率205000(MPa)、G:剪断剛性率79000(MPa)、
Mp:コラムの全塑性モーメントで下式による。
When the load (moment) decreases by 5% from the maximum value due to local buckling or brittle fracture, the specimen is considered to be fractured, and the total plastic rotation angle (cumulative plastic rotation angle Σθpl) of the specimen is calculated and tested. The cumulative plastic deformation magnification η was obtained as an index of the plastic deformability of the body, and when it was 30 or more, it was assumed that the structural member was excellent in earthquake resistance (plastic deformation performance). Η is calculated from the following equation.
η = Σθpl / θp
However, θp = (Pp / 2) L 2 / (3 · E · I) + Pp / 2 / (G · Aw)
Here, Pp: Total plastic load (N) = Mp / L,
L: Column cantilever length (distance from diaphragm to column end support point, 3250 mm)
E: Young's modulus 205000 (MPa), G: shear rigidity 79000 (MPa),
Mp: Total plastic moment of the column.

I:コラムの断面2次モーメント、σy:鋼材の降伏強度(MPa)
ここでIは下式による。
I: secondary moment of section of column, σy: yield strength of steel (MPa)
Here, I is according to the following equation.

D:コラム径(mm)、t:コラム板厚(mm)、
r:コラム角部内面の曲げ半径、R=r+t
Aw:剪断面積(mm)で下式による。
D: Column diameter (mm), t: Column plate thickness (mm),
r: bending radius of the inner surface of the column corner, R = r + t
Aw: Shear area (mm 2 ) according to the following formula.

表3に(1)から(5)の試験結果を示す。本発明例(厚鋼板No.2、4、5、10、11、15、16、17)はいずれも、降伏強さYS:385MPa以上、引張強さTS:550MPa以上、降伏比YR:75%以下を有し、さらに表層部および板厚方向1/4t位置でのvTrsが−40℃以下を満足する、高強度、高靭性の非調質低降伏比高張力厚鋼板である。   Table 3 shows the test results of (1) to (5). In all of the inventive examples (thick steel plates No. 2, 4, 5, 10, 11, 15, 16, 17), the yield strength YS: 385 MPa or more, the tensile strength TS: 550 MPa or more, the yield ratio YR: 75% A high-strength, high-toughness, non-tempered, low-yield-ratio, high-tensile steel plate having the following:

さらに、本発明例はいずれも、表層部の平均硬さが225HV以下で、表層部と板厚中央部との硬度差が60HV以下となる板厚方向硬さ分布を有し、冷間曲げを施しプレスコラムに加工し、プレスコラム−ダイアフラム接合部構造部材を構成した場合、プレスコラム−ダイアフラム接合部の3点曲げ試験における累積塑性変形倍率が30以上であり、耐震性能(塑性変形性能)に優れた、構造部材とすることができる。   Further, all of the examples of the present invention have a thickness distribution in the thickness direction in which the average hardness of the surface layer portion is 225 HV or less, and the hardness difference between the surface layer portion and the thickness center portion is 60 HV or less, and cold bending is performed. When a press column-diaphragm joint structural member is formed by processing into a press column, the cumulative plastic deformation ratio in the three-point bending test of the press column-diaphragm joint is 30 or more, and the seismic performance (plastic deformation performance) is improved. It can be set as the excellent structural member.

一方、成分組成および/またはミクロ組織が本発明の範囲外となる比較例(厚鋼板No.1、3、6〜9、12〜14、18〜24)は、強度、降伏比、靭性が不足しているか、冷間加工後の表層部の延性、靭性が低下し、構造部材としての累積塑性変形倍率ηが低くなっている。   On the other hand, comparative examples (thick steel plates No. 1, 3, 6-9, 12-14, 18-24) whose component composition and / or microstructure are outside the scope of the present invention are insufficient in strength, yield ratio, and toughness. In other words, the ductility and toughness of the surface layer portion after cold working are reduced, and the cumulative plastic deformation ratio η as a structural member is low.

表1に示す鋼No.A〜No.Eの組成を有する鋼素材を用いて、冷却工程のうち、第一段冷却を、冷却停止とその後の復熱を1回以上繰り返す加速冷却とし、第二段冷却を、冷却停止とその後の復熱を1回以上繰り返す加速冷却として、板厚:40mmの厚鋼板を製造した。得られた厚鋼板について、実施例1と同様の条件で、組織観察、硬さ試験、引張試験、衝撃試験およびコラム曲げ試験を実施した。   Steel No. shown in Table 1 A-No. Using the steel material having the composition E, in the cooling process, the first stage cooling is accelerated cooling in which the cooling stop and the subsequent recuperation are repeated one or more times, and the second stage cooling is performed in the cooling stop and the subsequent recovery. A plate thickness: 40 mm thick steel plate was manufactured as accelerated cooling in which heat was repeated once or more. The obtained thick steel plate was subjected to structure observation, hardness test, tensile test, impact test and column bending test under the same conditions as in Example 1.

表4に製造条件を、表5に試験結果を示す。本発明例(厚鋼板No.A1〜A5、A8)はいずれも、降伏強さYS:385MPa以上、引張強さTS:550MPa以上、降伏比YR:75%以下を有し、さらに表層部および板厚方向1/4t位置でのvTrsが−40℃以下を満足する、高強度、高靭性の非調質低降伏比高張力厚鋼板である。   Table 4 shows the manufacturing conditions, and Table 5 shows the test results. Examples of the present invention (thick steel plates No. A1 to A5, A8) all have a yield strength YS: 385 MPa or more, a tensile strength TS: 550 MPa or more, and a yield ratio YR: 75% or less. This is a high-strength, high-toughness, non-tempered, low yield ratio, high-tensile steel plate with a vTrs at a position of 1/4 t in the thickness direction of −40 ° C. or less.

本発明例はいずれも、表層部の平均硬さが225HV以下で、表層部と板厚中央部との硬度差が60HV以下となる板厚方向硬さ分布を有し、冷間曲げでプレスコラムに加工し、プレスコラム−ダイアフラム接合部構造部材を構成した場合、プレスコラム−ダイアフラム接合部の3点曲げ試験における累積塑性変形倍率が30以上の、耐震性能(塑性変形性能)に優れた、構造部材とすることができる。一方、本発明の範囲を外れる比較例(厚鋼板No.A6、A7)は、所望の引張強度または降伏強度、所望の降伏比が確保できていないか、所望の板厚方向硬さ分布が確保できておらず構造部材としての累積塑性変形倍率ηが低くなっている。   Each of the examples of the present invention has a thickness distribution in the thickness direction in which the average hardness of the surface layer portion is 225 HV or less, and the hardness difference between the surface layer portion and the plate thickness central portion is 60 HV or less. When a press column-diaphragm joint structural member is formed, the structure has excellent seismic performance (plastic deformation performance) with a cumulative plastic deformation ratio of 30 or more in a three-point bending test of the press column-diaphragm joint. It can be a member. On the other hand, in the comparative examples (thick steel plates No. A6, A7) outside the scope of the present invention, the desired tensile strength or yield strength, the desired yield ratio cannot be ensured, or the desired thickness distribution in the thickness direction is ensured. The cumulative plastic deformation magnification η as a structural member is low.

1 柱
2 通しダイアフラム
3 溶接部
4 裏当て金
1a プレスコラム
2a ダイアフラム(通しダイアフラム)
3a 4面BOX柱
4a、5a 溶接部
6 試験体
1 Column 2 Through Diaphragm 3 Welded Part 4 Backing Metal 1a Press Column 2a Diaphragm (Through Diaphragm)
3a 4-sided BOX column 4a, 5a Welded part 6 Specimen

Claims (7)

質量%で、C:0.05〜0.16%、Si:0.05〜0.45%、Mn:1.2〜1.8%、P:0.020%以下、S:0.005%以下、Al:0.05%以下、Ti:0.005〜0.020%、N:0.0040%以下、4.0≧Ti/N≧2.0、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%未満に制限し、さらに下記(1)式で定義されるCeqが、0.35〜0.48を満足し、残部Feおよび不可避的不純物からなる成分組成と、少なくとも鋼板の表層部において、平均結晶粒径が4.0〜18.0μmのフェライトと、パーライト、ベイナイトおよびマルテンサイトの1種または2種以上からなる硬質相からなり、フェライト面積率が45%〜70%のミクロ組織を有し、鋼板の表層部の平均硬さが225HV以下で、該表層部と板厚中央部との硬度差が60HV以下であることを特徴とする、冷間加工後の表層部の延性・靭性に優れる降伏強さ385MPa以上、引張強さ550MPa以上、降伏比75%以下である非調質低降伏比高張力厚鋼板。
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 (1)
(ここで、C、Mn、Cr、Mo、V、Cu、Niは各元素の含有量(質量%)で含有しない場合は0とする。)
In mass%, C: 0.05 to 0.16%, Si: 0.05 to 0.45%, Mn: 1.2 to 1.8%, P: 0.020% or less, S: 0.005 %: Al: 0.05% or less, Ti: 0.005-0.020%, N: 0.0040% or less, 4.0 ≧ Ti / N ≧ 2.0, and Nb, Mo as impurity elements Is limited to Nb: 0.004% or less and Mo: less than 0.04%, and Ceq defined by the following formula (1) satisfies 0.35 to 0.48, and the balance Fe and unavoidable Component composition consisting of impurities, at least in the surface layer portion of the steel sheet, consisting of a ferrite having an average crystal grain size of 4.0 to 18.0 μm and a hard phase consisting of one or more of pearlite, bainite and martensite, It has a microstructure with a ferrite area ratio of 45% to 70%, steel The yield strength is excellent in the ductility and toughness of the surface layer part after cold working, characterized in that the average hardness of the surface layer part is 225 HV or less and the hardness difference between the surface layer part and the plate thickness central part is 60 HV or less. A non-tempered low yield ratio high tensile steel plate having a thickness of 385 MPa or more, a tensile strength of 550 MPa or more, and a yield ratio of 75% or less.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
(Here, C, Mn, Cr, Mo, V, Cu, and Ni are set to 0 when not contained in the content (mass%) of each element.)
成分組成が、さらに質量%で、Cu:0.05〜0.50%、Ni:0.05〜0.80%、Cr:0.05〜0.60%、V:0.01〜0.05%、B:0.0003〜0.0030%の1種または2種以上を含有することを特徴とする請求項1に記載の非調質低降伏比高張力厚鋼板。 The component composition is further mass%, Cu: 0.05-0.50%, Ni: 0.05-0.80%, Cr: 0.05-0.60%, V: 0.01-0. The non-tempered low yield ratio high-tensile thick steel plate according to claim 1, containing one or more of 05% and B: 0.0003 to 0.0030%. 成分組成が、更に質量%で、Ca:0.0005〜0.0050%、REM:0.0010〜0.0050%の1種または2種を含有することを特徴とする請求項1または2記載の非調質低降伏比高張力厚鋼板。   The component composition further includes one or two of Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% in terms of mass%. Non-tempered low yield ratio high tensile steel plate. 更に、下記(2)式で定義されるACRが0.2〜0.8であることを特徴とする請求項3に記載の非調質低降伏比高張力厚鋼板。
ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S) (2)
(ここで、Ca、REM、O、Sは各元素の含有量(質量%)で含有しない場合は0とする。)
Furthermore, ACR defined by the following formula (2) is 0.2 to 0.8, the non-tempered low yield ratio high tensile thick steel plate according to claim 3.
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
(Here, Ca, REM, O, and S are set to 0 when not included in the content (mass%) of each element.)
成分組成が請求項1乃至4のいずれか一つに記載の鋼素材を1050〜1200℃に加熱後、表面温度で950℃以下の温度域での累積圧下量が30%以上で、圧延終了温度が表面温度で900℃以下Ar3変態点以上となる熱間圧延を行い、その後、第一段冷却として表面温度でAr3変態点以上の温度から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、表面温度が(Ar3変態点−100℃)以下550℃以上となるまで加速冷却し、冷却停止後復熱させ、第二段冷却として表面温度が(Ar3変態点−20℃)以下600℃以上、かつ、表面温度が極大値をとった時点から、下記(3)式を満たす時間t1(秒)以上、下記(4)式を満たす時間t2(秒)以内から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止後の復熱で表面温度が600℃以下400℃以上になる冷却停止温度まで加速冷却することを特徴とする非調質低降伏比高張力厚鋼板の製造方法。
After heating the steel material according to any one of claims 1 to 4 to 1050 to 1200 ° C, the cumulative reduction amount in the temperature range of 950 ° C or less at the surface temperature is 30% or more, and the rolling finish temperature Is subjected to hot rolling at a surface temperature of 900 ° C. or lower and an Ar3 transformation point or higher, and then, as the first stage cooling, from the temperature at the surface temperature of the Ar3 transformation point or higher, the average cooling at 1/4 position of the sheet thickness (t) At a rate of 2 ° C./s or more, accelerated cooling is performed until the surface temperature becomes (Ar 3 transformation point−100 ° C.) or less, 550 ° C. or more, recooling is performed after cooling is stopped, and the surface temperature is (Ar 3 transformation point− 20 ° C. or less, 600 ° C. or more, and from the time when the surface temperature takes the maximum value, from the time t1 (second) that satisfies the following formula (3) or more, and within the time t2 (second) that satisfies the following formula (4), Average cooling rate at 1/4 position of thickness (t) is 2 ° C / s or more The method of non-heat treated low yield ratio high-strength thick steel plate surface temperature at the recuperator after cooling stops characterized by accelerated cooling to a cooling stop temperature becomes 600 ° C. or less 400 ° C. or higher.
前記第一段冷却は、表面温度でAr3変態点以上の温度から冷却を開始し、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止温度が表面温度で550℃以上となる加速冷却を、複数回繰り返す冷却とし、該複数回の加速冷却において、冷却停止温度が表面温度で(Ar3変態点−100℃)以下550℃以上となる加速冷却を少なくとも1回含むことを特徴とする請求項5に記載の非調質低降伏比高張力厚鋼板の製造方法。   The first stage cooling starts from a temperature equal to or higher than the Ar3 transformation point at the surface temperature, is at an average cooling rate of 2 ° C./s or more at a 1/4 position of the plate thickness (t), and the cooling stop temperature is the surface temperature. The accelerated cooling at 550 ° C. or higher is repeated multiple times, and in the multiple times of accelerated cooling, the accelerated cooling at which the cooling stop temperature is not more than (Ar 3 transformation point−100 ° C.) at the surface temperature is 550 ° C. or higher at least once. The manufacturing method of the non-tempered low yield ratio high-tensile thick steel plate of Claim 5 characterized by the above-mentioned. 第二段冷却を、表面温度が(Ar3変態点−20℃)以下600℃以上、かつ、表面温度が極大値をとった時点から、(3)式を満たす時間t1(秒)以上、(4)式を満たす時間t2(秒)以内から、板厚(t)の1/4位置の平均冷却速度2℃/s以上で、冷却停止後復熱で表面温度が400℃以上となる冷却停止温度まで加速冷却する冷却を、複数回繰り返す冷却とし、前記複数回の加速冷却において、冷却停止後の復熱で表面温度が600℃以下400℃以上になる冷却停止温度まで冷却する加速冷却を最終冷却とすることを特徴とする請求項5または6に記載の非調質低降伏比高張力厚鋼板の製造方法。   The second stage cooling is performed at a time t1 (seconds) or more satisfying the expression (3) from the time when the surface temperature is (Ar3 transformation point −20 ° C.) or lower and 600 ° C. or higher and the surface temperature takes the maximum value. ) The cooling stop temperature at which the surface temperature becomes 400 ° C. or higher by reheating after cooling stop at an average cooling rate of 2 ° C./s or more at the 1/4 position of the plate thickness (t) from the time t2 (seconds) satisfying the formula The cooling to be accelerated cooling to the cooling is repeated a plurality of times, and the accelerated cooling to the cooling stop temperature at which the surface temperature becomes 600 ° C. or lower and 400 ° C. or higher by the recuperation after the cooling is stopped is the final cooling. The method for producing a non-tempered low yield ratio high tension thick steel plate according to claim 5 or 6.
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