JP2012107310A - Non-tempered low-yield-ratio high-tensile-strength steel plate and method for manufacturing the same - Google Patents

Non-tempered low-yield-ratio high-tensile-strength steel plate and method for manufacturing the same Download PDF

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JP2012107310A
JP2012107310A JP2011073363A JP2011073363A JP2012107310A JP 2012107310 A JP2012107310 A JP 2012107310A JP 2011073363 A JP2011073363 A JP 2011073363A JP 2011073363 A JP2011073363 A JP 2011073363A JP 2012107310 A JP2012107310 A JP 2012107310A
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JP5842359B2 (en
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Akio Omori
章夫 大森
Keiji Ueda
圭治 植田
Shinichi Suzuki
伸一 鈴木
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 non-tempered low-yield-ratio high-strength steel plate preferred for members of buildings and structures.SOLUTION: The non-tempered low-yield-ratio high-strength steel plate has such a composition that it contains 0.05-0.10% C, 1.2-1.8% Mn, 0.0010-0.0030% S, 0.005-0.020% Ti, and 0.0030-0.0060% N so that it satisfies a Ti/N of 2.0-4.0, it further contains at least one selected from among Cu, Ni, Cr, V, and B, it furthermore contains as impurities elements Nb and Mo provided that Nb is limited to 0.004% or less, and Mo is limited to 0.04% or less, and it satisfies a Ceq of 0.35-0.48; and has such a structure that at least the surface layer has ferrite and one or more selected from among pearlite, bainite, and martensite as a hard phase, and the average crystal grain diameter of the ferrite is 4.0-18.0 μm, furthermore the average hardness of the surface layer is 225 HV or less, and the difference in the hardness between the surface layer part and the center part of the plate thickness is 60 HV or less.

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、とくに例えば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、さらに大入熱溶接を施されて使用される使途に好適な、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 producing the same, and in particular, cold bending such as a circular steel pipe column or a square steel pipe column. The present invention relates to a non-tempered, low yield ratio, high-tensile steel plate and a method for producing the same, which are suitable for use that is applied after being subjected to high heat input welding.

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。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。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。   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. For this reason, building structures are obliged to use low yield ratio steel with a yield ratio (YR) of 80% or less. Furthermore, recently, with the increase in the height and the span of building structures, there are an increasing number of cases where 550 MPa class high-strength steel 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 has been 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 includes C: 0.02 to 0.04%, solute B: 0.0002 to 0.002%, and a composition that satisfies the range of 0.21 to 0.30% of the formula CEN related to the alloy element content, and bainite. 590MPa class non-tempered type low yield ratio high tensile strength steel sheet having a structure in which 0.8 to 2.5% by volume of island martensite is dispersed has been 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 it is necessary to further contain a large amount of alloying elements in order to ensure the desired high strength. 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%, containing P, S, Al, and N in an adjusted manner, and Mo and / or An island martensite (MA) containing W so as to satisfy a specific relational expression and having a Pcm of 0.22% or less and a structure in the central part of the plate thickness of ferrite as a main phase and 20% by volume or less. A high-tensile thick steel plate having a low yield ratio, which is a composite structure including a hard phase mainly comprising a phase), is described. With such a structure, a desired low yield ratio can be realized. In order to obtain such a structure, in the technique described in Patent Document 2, a steel material having the above composition is subjected to hot rolling with a rolling end temperature of 800 to 950 ° C. at a surface temperature of 0.5 to 50. It is said that it is preferable to sequentially perform a cooling process of accelerating cooling to a temperature range of 580 to 670 ° C. at an average cooling rate of ° C./s. However, the technique described in Patent Document 2 has a problem in that it requires expensive Mo and W to be contained, resulting in an increase in manufacturing cost.

また、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。   In Patent Document 3, steel containing C: 0.03 to 0.30%, Si: 0.05 to 0.60%, Mn: 0.50 to 2.5%, Al: 0.005 to 0.1% is heated, and the rolling end temperature is 900 ° C to The temperature is in the range of the Ar3 transformation point and the hot rolling to reduce the cumulative reduction in the temperature range to less than 30%, and air cooling after hot rolling, the surface temperature is (Ar3 transformation point -20 ° C) to (Ar3 transformation point) A method for producing a low yield ratio non-heat treated steel is described in which water cooling is started after reaching a temperature in the range of −80 ° C. and accelerated cooling is performed to stop cooling between 350 and 600 ° C.

また、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。   In Patent Document 4, steel containing C: 0.03 to 0.30%, Si: 0.05 to 0.60%, Mn: 0.50 to 2.5%, Al: 0.005 to 0.1% is heated, and the rolling end temperature is 900 ° C to The temperature is within the range of the Ar3 transformation point, the hot rolling is carried out so that the cumulative reduction rate in the temperature range is less than 30%, and air cooling is performed after the hot rolling. A method for producing a low yield ratio non-heat treated steel is described in which water cooling is started after reaching a temperature in the range of −80 ° C., accelerated cooling is performed until the temperature reaches 250 ° C. or less, and then tempering heat treatment is performed.

また、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: 0.003 to 0.030%, Ti: 0.005 to 0.020%, N : A steel slab containing 0.006% or less is hot-rolled so that the cumulative reduction at 900 ° C or less is 30% or more and the finishing temperature is Ar3 + 100 ° C or less Ar3 or more, and the steel sheet is (Ar3-20 ° C) to (Ar3- A method for producing a low yield ratio non-tempered steel is described in which air cooling to 100 ° C. is started, water cooling is started from this temperature, and cooling is stopped in the range of 400 to 550 ° C.

このように特許文献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。   As described above, 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, the steel slab is hot-rolled to complete the rolling above the Ar3 transformation point, and then is cooled to the austenite + ferrite two-phase temperature before starting accelerated cooling to produce 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。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 Patent Document 6, steel pieces containing 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: 0.001 to 0.010% are heated. Rough rolling with cumulative rolling reduction of 10-80% in the range up to 900 ° C and accelerated cooling of 2-40 ° C / s after rough rolling to (Ar3 transformation point + 50 ° C) to (Ar3 transformation point-50 ° C) And the austenite (γ) phase is supercooled, and finish rolling with a cumulative rolling reduction of 30 to 90% is completed at 650 ° C. or higher, and further accelerated cooling at 5 to 40 ° C./s is performed to 250 to 450 ° C., A method for producing a low yield ratio high strength steel is described. In the technique described in Patent Document 6, accelerated cooling is performed after rough rolling, the γ phase is subcooled to near Ar3 temperature, and then finish rolling is performed, so that fine ferrite is obtained 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 productivity while at the same time improving productivity. 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。 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 at least. Rapid cooling to the Ar1 point or lower at a speed, rolling starts or resumes after the surface layer reaches a temperature of the Ar3 point or lower, finishes rolling in the range of (Ac3-50 ° C) to Ac3, and then the Ac3 point or higher 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. ing. As a result, the structure of the front and back layers from 1 mm to 30% of the thickness of the steel sheet has a grain size more than three times the ferrite grain size inside the thickness, and the yield strength is greater than the yield strength of the inner thickness layer. It is said to be a low-yield-strength steel sheet with a surface layer lower by 5 kg / mm 2 or more.

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

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。
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と通しダイアフラムの接合部(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、2はダイアフラムで、4は当金である。
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.
Fig. 1 schematically shows the situation in which the joint (cross joint) between a column using a press column (cold-formed square steel pipe) or a circular steel pipe and a diaphragm is damaged when subjected to repeated tensile and compression deformation due to an earthquake. Indicate. 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. In addition, 2 is a diaphragm and 4 is a money.

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。
最近の建築構造物では、柱、梁等を、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を使用して構成することが多くなっている。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(鋼材)では、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。このため、このような冷間曲げ加工によって成形された鋼材を使用して、柱−梁接合部や柱−ダイアフラム接合部などのT継手や十字継手を形成すると、地震等による引張・圧縮繰り返し変形で表層に応力が集中した場合、早期に破断する危険性が高く、期待するような部材性能を発揮できない可能性がある。
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 are excellent.
In recent building structures, columns, beams, and the like are often configured using circular steel pipes and press columns formed by cold bending. In circular steel pipes and press columns (steel materials) formed by cold bending, the vicinity of the steel sheet surface layer is significantly hardened by cold bending, and the ductility and toughness near the surface layer are compared to the case where the steel sheet is used without being processed. Is in a lowered state. For this reason, when steel joints formed by such cold bending work are used to form T-joints and cross joints such as column-beam joints and column-diaphragm joints, repeated tensile and compression deformation due to earthquakes, etc. When stress concentrates on the surface layer, there is a high risk of breaking early, and there is a possibility that the expected member performance cannot be exhibited.

このような状況から、建築構造物部材用として、冷間加工を施された後においても、鋼板表層部の延性・靭性が優れた低降伏比高張力厚鋼板が要望されている。
しかし、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、鋼板表層付近での特性については全く考慮されておらず、上記した要望には対処できないという問題があった。
Under such circumstances, there is a demand for a low-yield-ratio high-tensile steel plate having excellent ductility and toughness of the steel sheet surface layer even after being cold worked for a building structure member.
However, all of the techniques described in Patent Documents 1 to 6 are mechanical properties (tensile properties) evaluated by full thickness tensile test pieces or round bar tensile test pieces at a thickness of 1/4 t or 1/2 t. , Ductility, toughness) is a technique for the purpose of obtaining desired characteristics, and the characteristics in the vicinity of the steel sheet surface layer are not considered at all, and there is a problem that the above-mentioned demand cannot be dealt with.

というのは、上記したような制御圧延や加速冷却(TMCP技術)を利用した非調質厚鋼板は、例えば図2の「曲げ加工前」のように、表層の硬さが最も高く,板厚中央の硬さが最も小さいという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。このような冷間曲げ加工を施され表層が著しく加工硬化した鋼材は、表層の延性・靭性が低下した状態となっている。このため、このような冷間曲げ加工を施された鋼板を、T継手や十字継手など表層に応力が集中する柱−梁接合部等の接合部に使用すると、全厚引張試験片または板厚1/4tや1/2t位置の丸棒引張試験片による引張特性が良好であっても、上記したように、接合部の表層から亀裂が発生し、接合部が早期破断する危険性が高くなることが予想される。   This is because non-tempered thick steel plates using controlled rolling and accelerated cooling (TMCP technology) as described above have the highest surface hardness, as shown in “Before bending” in FIG. It has a hardness distribution in the thickness direction where the hardness at the center is the smallest. 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. The steel material whose surface layer has been subjected to such cold bending and whose surface layer has been markedly work hardened is in a state where the surface layer has reduced ductility and toughness. For this reason, if a steel plate subjected to such cold bending is used for a joint such as a T-joint or a cross joint where a stress is concentrated on the surface layer such as a column-beam joint, Even if the tensile properties of the 1 / 4t or 1 / 2t round bar specimens are good, as described above, there is a higher risk of cracking from the surface layer of the joint and premature fracture of the joint. It is expected that.

このような問題に対し、例えば、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と考えられるが、特許文献7に記載された技術で製造された鋼板では、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恐れがあるという問題があった。   For such a problem, for example, 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 is improved. However, in the steel sheet manufactured by the technique described in Patent Document 7, the ferrite grains in the surface layer of the steel sheet are coarse, so it cannot be said that the toughness is sufficient, and the building structure is used as a member. When incorporated in an object, there is a problem that brittle fracture may occur from the member.

また、最近では、使用する鋼材の肉厚も厚くなり、プレスコラムや円形鋼管のシーム溶接に、例えば、サブマージアーク溶接等の、入熱:300kJ/cm以上の大入熱溶接を適用することが多くなっている。このため、建築構造物部材用として、優れた大入熱溶接部靭性を兼備した鋼材が要望されている。
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、さらに大入熱溶接部靭性にも優れた、降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:80%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。なお、ここでいう「優れた大入熱溶接部靭性」とは、入熱:300kJ/cm以上のサブマージアーク溶接の溶接ボンド部で、シャルピー衝撃試験による吸収エネルギーが、試験温度:0℃で70J以上である場合をいうものとする。また、ここでいう「厚鋼板」は、板厚:19mm以上、好ましくは25mm以上の鋼板をいうものとする。
Recently, the steel materials used have also become thicker, and it is possible to apply high heat input welding with a heat input of 300 kJ / cm or more, such as submerged arc welding, to seam welding of press columns and circular steel pipes. It is increasing. For this reason, steel materials having excellent large heat input weld toughness have been demanded for building structure members.
The present invention solves the above-mentioned problems of the prior art, uses a press column and a circular steel pipe without performing heat treatment such as quenching and tempering and normalizing and minimizing the alloy content. Yield that is suitable for building structure members, even after cold bending, with little increase in the hardness of the steel plate surface layer, excellent ductility and toughness of the steel plate surface layer, and excellent high heat input weld toughness An object of the present invention is to provide a non-tempered low yield ratio high tensile steel plate having a strength of 385 MPa or more, a tensile strength of 550 MPa or more, and a yield ratio of 80% or less, and a method for producing the same. The term “excellent large heat input weld toughness” as used herein refers to a weld bond part of submerged arc welding with a heat input of 300 kJ / cm or higher. The absorbed energy by the Charpy impact test is 70 J at a test temperature of 0 ° C. This is the case. In addition, the “thick steel plate” here refers to a steel plate having a thickness of 19 mm or more, preferably 25 mm or more.

本発明者らは、上記した目的を達成するために、非調質鋼板ではどうしても避けられない板厚方向に不均一な硬さ分布をある程度許容したうえで、建築構造物用部材として要求される所望の性能を確保するために、素材である鋼板が具備すべき性能について鋭意研究した。
冷間曲げ加工による塑性歪は、鋼板の表裏面で最大となり、板厚中央付近の中立点ではゼロとなる。このため、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。そこで、本発明者らは、冷間曲げ加工後の表層部で所望の延性・靭性を確保するためには、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ことが肝要であると考えた。その際、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。鋼板として所望の高強度を確保するためには、板厚中央部で一定以上の硬さ(強度)を確保することが必要となることに思い至った。
In order to achieve the above-mentioned object, the present inventors are required as a member for a building structure after allowing a non-uniform hardness distribution to some extent in the thickness direction, which is unavoidable with non-tempered steel sheets. In order to ensure the desired performance, we have intensively studied the performance that the steel sheet as the material should have.
The plastic strain due to cold bending is maximized on the front and back surfaces of the steel sheet, and is zero at the neutral point near the center of the sheet thickness. For this reason, the work hardening by cold bending becomes the most remarkable in the steel plate surface layer part. Therefore, in order to ensure the desired ductility and toughness in the surface layer portion after the cold bending process, the present inventors control the thickness direction hardness distribution before the cold bending process, We thought that it was important to reduce the hardness. At that time, if the hardness of the surface layer portion is lowered while the hardness of the central portion of the plate thickness is kept as it is, the strength at the full thickness of the steel plate is lowered. In order to secure a desired high strength as a steel plate, it has been thought that it is necessary to ensure a certain level of hardness (strength) at the center of the plate thickness.

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を、軟質相であるフェライト(好ましくは10面積%以上、より好ましくは30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あり、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ことも見出した。   In order to reduce the hardness in the vicinity of the surface layer part and achieve a low yield ratio, at least the microstructure of the surface layer part is ferrite (preferably 10 area% or more, more preferably 30 area% or more) that is a soft phase. ) To form a multiphase structure with the hard phase, and the ductility and toughness of the surface layer portion are desired by adjusting the average particle diameter of the ferrite in the surface layer portion within a desired appropriate range. It has also been found that it can be within the range.

そして、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ことを見出した。
(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。
And if it was a thick steel plate which satisfied following (1)-(4), it discovered that the deformation | transformation performance required as an object for building structure members could 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 plate has a multiphase structure composed of ferrite and a hard phase.
(2) The average hardness of the steel sheet surface layer 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.

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、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。また、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。   Here, the hard phase means a phase composed of one or more of pearlite, bainite, and martensite, and the steel sheet surface layer portion means a region of 1 to 5 mm in the thickness direction from the steel sheet front and back surfaces. The center of thickness refers to a region with a thickness center of ± 2 mm. In addition, the reason for limiting the structure and hardness of the steel sheet surface layer part was that the influence of the surface layer part, which is a 5 mm region in the sheet thickness direction from the steel sheet front surface or the back surface, was great for the destruction of the welded structure. Based on that. In addition, the outermost layer, which is less than 1 mm in the thickness direction from the front or back surface of the steel plate, is excluded because the outermost layer receives an extremely complicated thermal history due to rolling, accelerated cooling, etc. This is because it is often extremely difficult to control.

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に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 toughness improvement 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 heat-affected zone with high heat input welding. Nb and Mo are elements that improve the hardenability, greatly contribute to the formation of upper bainite containing 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, the present invention has found that it is necessary to strictly limit the content of Nb and Mo as impurities without adding Nb and Mo. Obtained.

本発明は、かかる知見に基づき、さらに検討を加えて完成されたものである。すなわち、本発明の要旨はつぎのとおりである。
(1)質量%で、C:0.05〜0.10%、Si:0.01〜0.45%、Mn:1.2〜1.8%、P:0.020%以下、S:0.0010〜0.0030%、Al:0.05%以下、Ti:0.005〜0.020%、N:0.0030〜0.0060%を含み、TiとNを、Ti含有量とN含有量との比、Ti/Nが2.0〜4.0を満足するように含有し、さらに、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種以上を含有し、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%以下に制限し、次(1)式
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 ‥‥(1)
(ここで、C、Mn、Cr、Mo、V、Cu、Ni:各元素の含有量(質量%))
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組成と、少なくとも、鋼板表面から板厚方向に1〜5mmの表層部がフェライトと、硬質相としてパーライト、ベイナイト、マルテンサイトのうち1種または2種以上からなり、前記フェライトの平均結晶粒径が4.0〜18.0μmである組織を有し、鋼板表面から板厚方向に1mm〜5mmの表層部の平均硬さが225HV以下で、該表層部と板厚中央位置を中心に±2mmの範囲である板厚中央部との硬度差が60HV以下である板厚方向硬さ分布を有し、冷間加工後の表層部延性・靭性に優れ、かつ大入熱溶接部靭性に優れることを特徴とする降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:80%以下である非調質低降伏比高張力厚鋼板。
The present invention has been completed based on such findings and further studies. That is, the gist of the present invention is as follows.
(1) By mass%, C: 0.05 to 0.10%, Si: 0.01 to 0.45%, Mn: 1.2 to 1.8%, P: 0.020% or less, S: 0.0010 to 0.0030%, Al: 0.05% or less, Ti: 0.005 -0.020%, N: 0.0030-0.0060% is included, Ti and N are contained so that ratio of Ti content and N content, Ti / N may satisfy 2.0-4.0, and Cu: 0.05- Containing one or more selected from 0.50%, Ni: 0.05-0.80%, Cr: 0.05-0.60%, V: 0.01-0.05%, B: 0.0003-0.0030%, and further an impurity element Nb and Mo are limited to Nb: 0.004% or less and Mo: 0.04% or less.
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
(Here, C, Mn, Cr, Mo, V, Cu, Ni: content of each element (mass%))
The carbon equivalent Ceq defined by the formula satisfies 0.35 to 0.48, the composition consisting of the remainder Fe and inevitable impurities, and at least the surface layer portion of 1 to 5 mm in the thickness direction from the steel plate surface is ferrite and pearlite as the hard phase , Bainite, martensite, or one or more of them, the ferrite has an average crystal grain size of 4.0 to 18.0 μm, and the average of the surface layer portion of 1 mm to 5 mm in the thickness direction from the steel sheet surface It has a hardness distribution in the plate thickness direction with a hardness difference of 60HV or less between the surface layer and the plate thickness center that is within ± 2mm centered on the plate thickness center position. Yield strength: 385MPa or more, tensile strength: 550MPa or more, yield ratio: 80% or less, characterized by excellent surface ductility and toughness afterwards, and high heat input weld toughness Yield ratio high tensile steel plate.

(2)(1)において、前記組成に加えてさらに、質量%で、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種を含有する組成とすることを特徴とする非調質低降伏比高張力厚鋼板。
(3)(2)において、前記Ca、REMを、次(2)式
ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S)‥‥ (2)
(ここで、Ca、REM、O、S:各元素の含有量(質量%))
で定義されるACRが0.2〜0.8を満足するように含有することを特徴とする非調質低降伏比高張力厚鋼板。
(2) In (1), in addition to the above composition, the composition further contains one or two kinds selected from Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% by mass%. Non-tempered low yield ratio high tensile thick steel sheet characterized by
(3) In (2), the Ca and REM are expressed by the following formula (2)
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
(Where Ca, REM, O, S: content of each element (mass%))
A non-tempered, low yield ratio, high-tensile steel plate characterized by containing an ACR defined by the above so as to satisfy 0.2 to 0.8.

(4)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、該厚鋼板に途中冷却停止を含む第一段冷却と第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、を施す非調質厚鋼板の製造方法において、前記鋼素材を、質量%で、C:0.05〜0.10%、Si:0.05〜0.45%、Mn:1.2〜1.8%、P:0.020%以下、S:0.0010〜0.0030%、Al:0.05%以下、Ti:0.005〜0.020%、N:0.0030〜0.0060%を含み、TiとNを、Ti含有量とN含有量との比、Ti/Nが2.0〜4.0を満足するように含有し、さらに、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種以上を含有し、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%以下に制限し、次(1)式
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 ‥‥(1)
(ここで、C、Mn、Cr、Mo、V、Cu、Ni:各元素の含有量(質量%))
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組成を有する鋼素材とし、前記熱間圧延の加熱温度を1050〜1200℃とし、前記熱間圧延を、表面温度で900℃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870℃以下Ar3変態点以上となる圧延とし、前記第一段冷却が、表面温度でAr3変態点以上の温度から冷却を開始し、板厚1/4t位置の平均冷却速度で3〜30℃/sの冷却速度で冷却し、表面温度が(Ar3変態点−100℃)以下400℃以上で、加速冷却を停止する冷却とし、冷却停止後、複熱し、表面温度が(Ar3変態点+10℃)以下650℃以上、表面と板厚中央の温度差が80℃以下となる時点で、前記第二段冷却を開始し、該第二段冷却を、板厚1/4t位置の平均冷却速度で3℃/s以上の冷却速度で、該第二段冷却を停止した後の復熱で表面温度が600℃以下になるような冷却停止温度まで加速冷却する冷却とすることを特徴とする降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:80%以下を有し、冷間加工後の表層部延性・靭性に優れ、かつ大入熱溶接部靭性に優れた非調質低降伏比高張力厚鋼板の製造方法。
(4) A steel material is subjected to hot rolling to form a thick steel plate, and following the rolling step, the steel plate is a two-stage cooling process comprising a first stage cooling and a second stage cooling including an intermediate cooling stop. In the manufacturing method of the non-heat treated thick steel sheet to be subjected to a cooling step for performing accelerated cooling, the steel material in mass%, C: 0.05 to 0.10%, Si: 0.05 to 0.45%, Mn: 1.2 to 1.8%, P: 0.020% or less, S: 0.0010 to 0.0030%, Al: 0.05% or less, Ti: 0.005 to 0.020%, N: 0.0030 to 0.0060%, Ti and N, the ratio of Ti content to N content , Ti / N is contained so as to satisfy 2.0 to 4.0, and 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 Containing one or more selected from 0.0030%, further limiting Nb and Mo as impurity elements to Nb: 0.004% or less and Mo: 0.04% or less, the following formula (1)
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
(Here, C, Mn, Cr, Mo, V, Cu, Ni: content of each element (mass%))
The carbon equivalent Ceq defined by the above is a steel material having a composition consisting of the balance Fe and inevitable impurities satisfying 0.35-0.48, the heating temperature of the hot rolling is 1050-1200 ° C., and the hot rolling is performed. The rolling reduction is such that the cumulative reduction amount in the temperature range of 900 ° C. or less at the surface temperature is 30% or more and the rolling end temperature is 870 ° C. or less and the Ar 3 transformation point or more at the surface temperature, and the first stage cooling is performed at the surface temperature Cooling is started from the temperature above the Ar 3 transformation point, cooling is performed at a cooling rate of 3 to 30 ° C / s at an average cooling rate at a thickness of 1 / 4t, and the surface temperature is (Ar 3 transformation point – 100 ° C) Accelerated cooling is stopped at 400 ° C or higher, cooling is stopped, double heat is applied after cooling stops, surface temperature is (Ar 3 transformation point + 10 ° C) or lower, 650 ° C or higher, and temperature difference between surface and thickness center is 80 ° C or lower. At this point, the second stage cooling is started, and the second stage cooling is performed at a cooling rate of 3 ° C./s or more at an average cooling rate at a thickness of 1/4 t. Yield strength: 385MPa or more, tensile strength: 550MPa or more, yield, characterized by accelerated cooling to a cooling stop temperature such that the surface temperature becomes 600 ° C or less by reheating after stage cooling is stopped Ratio: 80% or less, a method for producing a non-tempered low yield ratio high-tensile thick steel plate having excellent surface layer ductility and toughness after cold working and excellent high heat input weld toughness.

(5)(4)において、前記第一段冷却に代えて、第一段冷却を、表面温度でAr3変態点以上の温度から冷却を開始し、板厚1/4t位置の平均冷却速度で2℃/s以上の冷却速度で、冷却停止温度が表面温度で400℃以上となる加速冷却を、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し、前記複数回の加速冷却が、冷却停止温度が表面温度で(Ar3変態点−100℃)以下400℃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非調質低降伏比高張力厚鋼板の製造方法。   (5) In (4), instead of the first-stage cooling, the first-stage cooling is started from the surface temperature at a temperature equal to or higher than the Ar3 transformation point, and the average cooling rate at the thickness of 1/4 t is 2 Accelerated cooling where the cooling stop temperature is 400 ° C. or more at the surface temperature at a cooling rate of ℃ / s or more is assumed to be repeated cooling multiple times with the cooling stop and subsequent recuperation sandwiched, A method for producing a non-tempered, low yield ratio, high-tensile steel plate, comprising at least one accelerated cooling at which the cooling stop temperature is 400 ° C. or higher at a surface temperature (Ar3 transformation point−100 ° C.).

(6)(4)または(5)において、前記第二段冷却に代えて、第二段冷却を、板厚1/4t位置の平均冷却速度で2℃/s以上の冷却速度で、冷却停止とその後の復熱とを挟んで、加速冷却を複数回繰り返す冷却とし、前記複数回の加速冷却のうち、冷却停止後の復熱で表面温度が600℃以下になるような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非調質低降伏比高張力厚鋼板の製造方法。   (6) In (4) or (5), in place of the second-stage cooling, the second-stage cooling is stopped at an average cooling rate at a thickness of 1/4 t at a cooling rate of 2 ° C./s or more. And the subsequent recuperation, the accelerated cooling is repeated multiple times. Of the multiple accelerated cooling, cooling to a cooling stop temperature such that the surface temperature is 600 ° C. or less due to the recuperation after the cooling is stopped. A method for producing a non-tempered low yield ratio high-tensile thick steel plate, characterized in that accelerated cooling is performed as final cooling.

(7)(4)ないし(6)のいずれかにおいて、前記冷却工程に引続き、400℃以上700℃以下の温度で焼戻しを行う焼戻工程を施す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。
(8)(4)ないし(7)のいずれかにおいて、前記鋼素材の組成に加えて、さらに、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種を含有する組成と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。
(7) In any one of (4) to (6), following the cooling step, a tempering step of tempering at a temperature of 400 ° C. or higher and 700 ° C. or lower is performed. A method for producing a tension thick steel plate.
(8) In any one of (4) to (7), in addition to the composition of the steel material, one or two selected from Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% A method for producing a non-tempered, low yield ratio, high-tensile thick steel plate, characterized by comprising a composition containing:

(9)(8)において、前記Ca、REMを、次(2)式
ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S)‥‥(2)
(ここで、Ca、REM、O、S:各元素の含有量(質量%))
で定義されるACRが0.2〜0.8を満足するように含有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。
(9) In (8), the Ca and REM are expressed by the following formula (2)
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
(Where Ca, REM, O, S: content of each element (mass%))
A method for producing a non-tempered, low yield ratio, high-tensile thick steel plate, wherein the ACR is defined so as to satisfy 0.2 to 0.8.

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、靭性に優れるとともに、大入熱溶接熱影響部靭性にも優れ、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:80%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 the hardness of the steel sheet surface layer part, the ductility and the toughness of the steel sheet surface layer part are excellent, and the high heat input heat affected zone toughness is also excellent. Suitable for materials, heat treatment is applied to non-tempered low yield ratio high strength thick steel plate with high strength of yield strength: 385MPa or higher, tensile strength: 550MPa or higher and yield ratio: low yield ratio of 80% or less It can be produced without containing a large amount of alloy, 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模式的に示す説明図である。It is explanatory drawing which shows typically 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模式的に示す説明図である。It is explanatory drawing which shows typically plate | board thickness direction hardness distribution of a non-tempered thick steel plate. 実施例におけるシャルピー衝撃試験片の採取位置を示す説明図である。It is explanatory drawing which shows the collection position of the Charpy impact test piece in an Example. 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)方法の概略を模式的に示す説明図である。It is explanatory drawing which shows typically the outline of the three-point bending test (column bending test) method of a press column-diaphragm junction part. 三点曲げ試験(コラム曲げ試験)の試験体における溶接部近傍を模式的に示す説明図である。It is explanatory drawing which shows typically the welding part vicinity in the test body of a three-point bending test (column bending test). 三点曲げ試験(コラム曲げ試験)の試験要領を模式的に示す説明図である。It is explanatory drawing which shows typically the test point of a three-point bending test (column bending test). 三点曲げ試験(コラム曲げ試験)における荷重−変形量ヒステリシス曲線の例を模式的に示すグラフである。It is a graph which shows typically an example of a load-deformation amount hysteresis curve in a three-point bending test (column bending test). 本発明における冷却工程の冷却条件の一例を模式的に示す説明図である。It is explanatory drawing which shows typically an example of the cooling conditions of the cooling process in this invention.

まず、本発明厚鋼板の組成限定理由について説明する。なお、以下、とくに断わらない限り、質量%は単に%で記す。
C:0.05〜0.10%
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、0.10%を超える含有は、大入熱溶接熱影響部の靭性が顕著に低下する。このため、Cは0.05〜0.10%の範囲に限定した。なお、好ましくは0.06〜0.09%である。
First, the reasons for limiting the composition of the steel plate of the present invention will be described. Hereinafter, unless otherwise specified, mass% is simply expressed as%.
C: 0.05-0.10%
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 0.10%, the toughness of the heat-affected zone with high heat input is significantly reduced. For this reason, C was limited to the range of 0.05 to 0.10%. In addition, Preferably it is 0.06 to 0.09%.

Si:0.01〜0.45%
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0.01%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)靱性を顕著に低下させる。このため、Siは0.01〜0.45%の範囲に限定した。なお、好ましくは、0.05〜0.35%である。
Si: 0.01-0.45%
Si acts as a deoxidizer and dissolves in the steel to increase the strength of the steel. In order to acquire such an effect, 0.01% 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 weld heat affected zone (HAZ) toughness is markedly lowered. For this reason, Si was limited to the range of 0.01 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: 1.2-1.8%
Mn is an element that has the effect of increasing the strength of the steel by solid solution, and is inexpensive, and in the present invention that minimizes the inclusion of other expensive alloy elements, the desired high strength (tensile strength: 550 MPa) In order to ensure the above), a content of 1.2% or more is required. On the other hand, if the content exceeds 1.8%, the toughness of the base metal and the HAZ toughness are significantly reduced. For this reason, Mn was limited to the range of 1.2 to 1.8%. 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 having an action of increasing the strength of steel, but is an element that lowers toughness, particularly toughness of a welded portion. In the present invention, it is desirable to reduce as much as possible, but excessive reduction However, it is preferable to make the refining cost about 0.005% or more because it raises the refining cost and is economically disadvantageous. 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:0.0010〜0.0030%
Sは、鋼中ではMnS等の硫化物系介在物として存在して、オーステナイト(γ)→フェライト(α)変態の核として作用し、溶接部靭性を向上させる作用を有する。このような効果は、0.0010%以上の含有で認められる。一方、0.0030%を超える含有は、鋳片中央偏析部などに多量のMnSが生成し、靭性を低下させるとともに、鋳片等における欠陥を発生しやすくする。このため、Sは0.0010〜0.0030%の範囲に限定した。なお、好ましくは0.0010〜0.0025%である。
S: 0.0010 to 0.0030%
S is present in the steel as sulfide inclusions such as MnS and acts as a nucleus of the austenite (γ) → ferrite (α) transformation and has an effect of improving the toughness of the weld. Such an effect is recognized when the content is 0.0010% or more. On the other hand, if the content exceeds 0.0030%, a large amount of MnS is generated in the center segregation portion of the slab, and the toughness is lowered and defects in the slab and the like are easily generated. For this reason, S was limited to 0.0010 to 0.0030% of range. In addition, Preferably it is 0.0010 to 0.0025%.

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 the 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 metal is lowered and mixed with the weld metal during welding to lower the weld metal toughness. Let For this reason, Al was limited to 0.05% or less. In addition, Preferably it is 0.010 to 0.045%.

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 with N, and precipitates as TiN during solidification, thereby reducing solute N in the steel and reducing 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 the range of 0.005-0.020%. In addition, Preferably it is 0.007 to 0.015%.

N:0.0030〜0.0060%
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が、TiNとして析出すると、とくに大入熱溶接熱影響部のγ粒の粗大化を抑制し、溶接熱影響部靭性を向上させる効果を有する。このような効果を得るためには、0.0030%以上含有することが必要であるが、0.0060%を超える含有は、靭性の劣化が著しくなる。このため、Nは0.0030〜0.0060%に限定した。
N: 0.0030-0.0060%
N, when dissolved in steel, causes strain aging after cold working and deteriorates toughness. However, when precipitated as TiN, it suppresses the coarsening of γ grains, especially in the heat-affected zone of high heat input welding. And has the effect of improving the weld heat affected zone toughness. In order to obtain such an effect, it is necessary to contain 0.0030% or more, but if it exceeds 0.0060%, the toughness deteriorates remarkably. For this reason, N was limited to 0.0030-0.0060%.

Ti/N:2.0〜4.0
本発明では、N含有量に見合う量のTiを含有させ、固溶NをTiNとして固定する。このため、Ti含有量とN含有量との比、Ti/Nが2.0以上を満足するように、Ti含有量を調整する。Ti/Nが2.0未満では、N含有量に比べてTi含有量が少なすぎ、多くのNが固溶Nとして残存する。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は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。
Ti / N: 2.0 to 4.0
In the present invention, an amount of Ti commensurate with the N content is contained, and the solid solution N is fixed as TiN. For this reason, Ti content is adjusted so that ratio of Ti content and N content, Ti / N may satisfy 2.0 or more. When Ti / N is less than 2.0, the Ti content is too small compared to the N content, and much N remains as solid solution N. On the other hand, when Ti / N is larger than 4.0, TiN particles are coarsened and a desired effect cannot be secured. 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.

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、B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1種または2種以上を含有する。
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、0.05%以上含有するが、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%の範囲に限定した。なお、より好ましくは0.10〜0.40%である。
One or more selected from 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 0.0030%
Cu, Ni, Cr, V, and B are all elements that have an action of increasing the strength of steel, and optionally contain one or more.
Cu contributes to increasing the strength of thick steel plates by increasing the strength of steel plates through solid solution strengthening and hardenability improvement. In order to obtain such an effect, the content is 0.05% or more, but the content exceeding 0.50% causes an increase in alloy cost and deterioration of surface properties due to hot brittleness. For this reason, when it contained, Cu was limited 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、0.05%以上含有するが、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。なお、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 almost no deterioration in toughness, and has little adverse effect on HAZ toughness, and is an element useful for increasing the strength of thick steel sheets. In order to acquire such an effect, it contains 0.05% or more, but if it contains more than 0.80%, Ni is an expensive element, which causes an increase in alloy costs. For this reason, when it contains, Ni is limited to 0.05 to 0.80%. In addition, 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。なお、好ましくは0.10〜0.50%である。
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、0.01%以上含有するが、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0.01〜0.06%の範囲に限定する。なお、好ましくは0.02〜0.04%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 contains 0.05% or more, but the content exceeding 0.60% causes an increase in alloy cost. For this reason, when it contains, Cr is limited to 0.05 to 0.60% of range. In addition, Preferably it is 0.10 to 0.50%.
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 contains 0.01% or more, but inclusion exceeding 0.05% reduces the toughness of a base material and HAZ. For this reason, when it contains, V is limited to 0.01 to 0.06% of range. In addition, Preferably it is 0.02 to 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%以上含有するが、0.0030%を超える含有は、母材やHAZの靭性を劣化させる。このため、含有する場合には、Bは0.0003〜0.0030%の範囲に限定した。なお、好ましくは0.0006〜0.0020%である。
さらに、本発明では、不純物としてのNb、Moの含有量を所定値以下に制限する。
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 0.0003% or more, but the content exceeding 0.0030% deteriorates the toughness of the base material and the HAZ. For this reason, when contained, B is limited to a range of 0.0003 to 0.0030%. In addition, Preferably it is 0.0006 to 0.0020%.
Furthermore, in the present invention, the contents of Nb and Mo as impurities are limited to a predetermined value or less.

Nb:0.004%以下、Mo:0.04%以下
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。しかし、不可避的不純物として含有される場合があるが、そのような場合でも、Nb:0.004%以下、Mo:0.04%以下に限定する。不可避的不純物として、Nbが0.004%を超えて、および/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%以下を満足させるためには、Nbや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が肝要となる。
さらに本発明では、上記した成分を上記した範囲で、かつ、(1)式で定義される炭素当量Ceqが、0.35〜0.48の範囲を満足するように、調整して含有する。
炭素当量Ceqは、次(1)式
Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 ‥‥(1)
(ここで、C、Mn、Cr、Mo、V、Cu、Ni:各元素の含有量(質量%))
で定義される。
Nb: 0.004% or less, Mo: 0.04% or less
Nb and Mo are elements that improve the hardenability, make it easy to generate upper bainite containing island martensite, and lower the toughness of the heat-affected zone with high heat input welding. For this reason, Nb and Mo are not added in the present invention. However, it may be contained as an unavoidable impurity, but even in such a case, it is limited to Nb: 0.004% or less and Mo: 0.04% or less. If Nb exceeds 0.004% and / or Mo exceeds 0.04% as an unavoidable impurity, the toughness of the heat input zone affected by high heat input welding decreases. In order to satisfy Nb: 0.004% or less and Mo: 0.04% or less, it is important to use raw materials with a low content of Nb and Mo, smelting furnace refractories, and the like.
Further, in the present invention, the above-described components are adjusted and contained so that the carbon equivalent Ceq defined by the formula (1) satisfies the range of 0.35 to 0.48 within the above-described range.
The carbon equivalent Ceq is the following formula (1)
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
(Here, C, Mn, Cr, Mo, V, Cu, Ni: content of each element (mass%))
Defined by

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。
上記した成分が基本の成分であるが、これら基本成分に加えて、必要に応じて、選択元素として、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。
If Ceq is less than 0.35, the desired base metal strength cannot be secured, and softening of the weld heat affected zone cannot be suppressed within the desired allowable limit. On the other hand, when Ceq exceeds 0.48, the weldability decreases, and the base metal toughness and HAZ toughness decrease. For this reason, Ceq was limited to the range of 0.35-0.48. In addition, Preferably it is 0.36-0.46.
The above-mentioned components are basic components. In addition to these basic components, one or two selected from Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% as optional elements as necessary. Can contain seeds.

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
Ca、REMはいずれも、硫化物の形態制御を介して母材の靭性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では少なくとも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。
One or two selected from Ca: 0.0005 to 0.0050%, REM: 0.0010 to 0.0050%
Both Ca and REM contribute to the improvement of the toughness and ductility of the base material through the control of sulfide morphology. In addition, when fine sulfide particles are dispersed in steel, they contribute to the improvement of HAZ toughness by acting as ferrite transformation nuclei. In order to obtain these effects, it is preferable to contain at least 0.0005% for Ca and at least 0.010% for REM, but if both Ca and REM contain more than 0.0050%, excessive inclusions are generated, Conversely, the toughness may decrease. 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、次(2)式
ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S) ‥‥ (2)
(ここで、Ca、REM、O、S:各元素の含有量(質量%))
で定義されるACRが、0.2〜0.8の範囲を満足するように、調整することが好ましい。ACRをこの範囲に調整すると、鋳造時に鋼中に微細なCaおよび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表面にMnSが析出する。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。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あるいは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。
When Ca and REM are contained, the following formula (2) is used in order to ensure the action of controlling the form of sulfide.
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
(Where Ca, REM, O, S: content of each element (mass%))
It is preferable to adjust so that the ACR defined by the above satisfies the range of 0.2 to 0.8. When the ACR is adjusted within this range, fine Ca and REM sulfates (oxysulfides) are dispersed in the steel during casting, and MnS precipitates on the inclusion surface during cooling after welding. 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. If the ACR is less than 0.2, the amount of Ca and REM is insufficient, and desired composite inclusions / precipitates cannot be generated, and MnS harmful to the base metal and weld heat affected zone toughness increases. On the other hand, when ACR exceeds 0.8, most of S is taken into the inclusions of Ca or REM, MnS precipitated on the inclusion surface is insufficient, and the inclusions do not sufficiently function as intragranular ferrite formation sites. For this reason, it is preferable to limit ACR to the range of 0.2-0.8.

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。
本発明厚鋼板は、上記した組成を有し、さらに、少なくとも鋼板表面から板厚方向に1〜5mmの領域である表層部が、フェライトと、硬質相としてパーライト、ベイナイト、マルテンサイトのうち1種または2種以上からなり、表層部における該フェライトの平均結晶粒径が4.0〜18.0μmであるミクロ組織を有する。
The balance other than the components described above is composed of Fe and inevitable impurities. As an inevitable impurity, O: 0.005% or less is acceptable.
The steel plate according to the present invention has the above-described composition, and at least a surface layer portion of 1 to 5 mm in the plate thickness direction from the steel plate surface is one of ferrite and pearlite, bainite, martensite as a hard phase. Or it consists of 2 or more types, and has the microstructure whose average crystal grain diameter of this ferrite in a surface layer part is 4.0-18.0 micrometers.

本発明厚鋼板では、降伏比:80%以下の低降伏比と引張強さ:550MPa以上の強度とを兼備させるために、鋼板表面から板厚方向に1mm未満の領域である最表層を除く、少なくとも鋼板表面から板厚方向に1〜5mmの領域である表層部を、好ましくは板厚方向全域を、軟質相であるフェライトと、硬質相からなる複相組織とする。
少なくとも表層部を、好ましくは最表層を除く板厚全域を、軟質相であるフェライトと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。とくに、優れた延性と低降伏比とを両立させるためには、軟質相であるフェライトは面積率で少なくとも10%以上含有することが好ましく、より好ましくは30%以上である。軟質相であるフェライトが面積率で10%未満では、とくに低降伏比を実現することができなくなる。なお、所望の高強度(引張強さ:550MPa以上)を確保するためには、フェライトは70%以下とすることが望ましい。
In the steel plate of the present invention, in order to combine the yield ratio: low yield ratio of 80% or less and the tensile strength: strength of 550 MPa or more, the outermost layer which is an area less than 1 mm in the thickness direction from the steel sheet surface is excluded. At least the surface layer portion that is 1 to 5 mm in the plate thickness direction from the surface of the steel plate, preferably the entire region in the plate thickness direction, has a multiphase structure composed of a soft phase ferrite and a hard phase.
At least the surface layer portion, preferably the entire plate thickness excluding the outermost layer, is made of a multiphase structure combining a soft phase ferrite and a hard phase, so that excellent ductility, desired high strength, and a low yield ratio are obtained. Both can be achieved. In particular, in order to achieve both excellent ductility and a low yield ratio, the soft phase ferrite is preferably contained in an area ratio of at least 10%, more preferably 30% or more. If the ferrite, which is a soft phase, is less than 10% in area ratio, a low yield ratio cannot be achieved. In order to secure a desired high strength (tensile strength: 550 MPa or more), the ferrite content is desirably 70% or less.

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。フェライトとこれら硬質相とを組み合わせた複相組織とすることにより,高強度と優れた延性を両立することができる。なお、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。
さらに、本発明厚鋼板では、鋼板表面から板厚方向に1mm〜5mmの表層部におけるフェライトの平均結晶粒径を4.0〜18.0μmとする。
The hard phase is one or more selected from pearlite, bainite, and martensite. High strength and excellent ductility can be achieved by using a multiphase structure combining ferrite and these hard phases. In addition, the kind of hard phase and those fractions can be suitably selected by desired intensity | strength and toughness, and also a chemical component and board thickness.
Furthermore, in the steel plate of the present invention, the average crystal grain size of ferrite in the surface layer portion of 1 mm to 5 mm in the plate thickness direction from the steel plate surface is set to 4.0 to 18.0 μm.

鋼板表層部における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。鋼板表層部のフェライト粒径が4.0μm未満では、降伏比が急激に上昇し、均一伸びが低下する。このため、鋼板表層部の塑性変形能が大きく低下する。鋼板表層部の塑性変形能の低下は、地震による変形時に溶接止端部などから亀裂を発生しやすくなる。また、鋼板表層部のフェライト粒径が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。このため、鋼板表層部のフェライトの平均結晶粒径は4.0〜18.0μmの範囲に限定した。なお、好ましくは6.0〜15.0μmである。   The average crystal grain size of ferrite in the surface layer portion of the steel plate greatly affects toughness, yield ratio and elongation. If the ferrite grain size of the steel sheet surface layer is less than 4.0 μm, the yield ratio increases rapidly and the uniform elongation decreases. For this reason, the plastic deformability of a steel plate surface layer part falls significantly. The decrease in the plastic deformability of the steel sheet surface layer part is liable to generate a crack from the weld toe during deformation due to an earthquake. Further, when the ferrite grain size of the steel sheet surface layer exceeds 18.0 μm and becomes coarse, the toughness is lowered and brittle fracture is likely to occur. For this reason, the average grain size of ferrite in the steel sheet surface layer portion is limited to a range of 4.0 to 18.0 μm. In addition, Preferably it is 6.0-15.0 micrometers.

さらに、本発明厚鋼板では、鋼板表面から板厚方向に1mm〜5mmの表層部の平均硬さが225HV以下で、該表層部と板厚中央位置を中心に±2mmの範囲である板厚中央部との硬度差が60HV以下である板厚方向硬さ分布を有する。
上記したような板厚方向硬さ分布とすることにより、冷間曲げ加工等を施される建築構造部材用として、必要な性能(例えば、塑性変形能、脆性破壊防止能等)を確保できる。
Furthermore, in the steel plate of the present invention, the average hardness of the surface layer portion of 1 mm to 5 mm in the plate thickness direction from the steel plate surface is 225 HV or less, and the plate thickness center is in the range of ± 2 mm centering on the surface layer portion and the plate thickness center position. The thickness difference in hardness in the plate thickness direction is 60HV or less.
By setting the thickness distribution in the plate thickness direction as described above, necessary performance (for example, plastic deformability, brittle fracture prevention capability, etc.) can be secured for a building structural member subjected to cold bending or the like.

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から亀裂を発生しやすくなる。なお、表層部以外に硬さが225HVを超える領域が存在したとしても、建築構造物の柱−梁接合部などの部材では、地震による変形時には溶接止端部など鋼板表面に大きな歪が集中するため、部材としての変形性能を大きく損ねることはない。   If the hardness of the steel plate surface layer exceeds 225 HV, after the cold bending process, the hardness of the steel plate surface layer portion will further increase, and the toughness and ductility of the steel plate surface layer portion will be significantly reduced. In members such as beam joints (T joints and cross joints), cracks are likely to occur from the weld toes when deformed by an earthquake. In addition, even if there is a region where the hardness exceeds 225HV other than the surface layer part, in a member such as a column-beam joint part of a building structure, a large strain concentrates on the steel plate surface such as a weld toe at the time of deformation due to an earthquake. Therefore, the deformation performance as a member is not greatly impaired.

つぎに、上記した本発明厚鋼板の好ましい製造方法について説明する。
本発明厚鋼板の製造方法では、上記した組成の鋼素材に、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、該厚鋼板に途中冷却停止を含む第一段冷却と第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施す。
本発明で使用する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。
Next, a preferred method for producing the above-described thick steel plate of the present invention will be described.
In the method for producing a thick steel plate according to the present invention, the steel material having the above composition is subjected to hot rolling to obtain a thick steel plate, and following the rolling step, the first stage cooling including a halfway cooling stop to the thick steel plate, And a cooling step for performing two-stage accelerated cooling consisting of second-stage cooling.
The manufacturing method of the steel material used in the present invention is not particularly limited, and any conventional melting method and casting method can be applied. However, the molten steel having the above composition is converted into a converter, electric furnace, vacuum melting furnace. It is preferable to use a continuous casting method or the like to produce a steel material (slab) through a deoxidation process 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され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900℃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870℃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。
加熱温度:1050〜1200℃
加熱温度が1050℃未満では、得られる厚鋼板の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厚鋼板の靱性が低下したり、焼入性が増加しすぎて得られる厚鋼板の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。
The obtained steel material (slab) is first heated and subjected to a rolling process in which it is hot rolled into a thick steel plate. In the rolling process, the steel material is heated to a heating temperature of 1050-1200 ° C, the cumulative reduction in the temperature range of 900 ° C or less at the surface temperature is 30% or more, and the rolling end temperature is 870 ° C or less at the surface temperature Ar3 transformation Hot rolling at a point or more is performed to obtain a thick steel plate having a predetermined thickness.
Heating temperature: 1050 ~ 1200 ℃
If the heating temperature is less than 1050 ° C, the strength of the resulting thick steel plate tends to decrease, while if it exceeds 1200 ° C, the toughness of the thick steel plate obtained by coarsening the structure decreases or the hardenability increases too much. The surface hardness of the thick steel plate obtained in this way tends to increase. 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900℃以下の温度域での累積圧下量:30%以上
本発明では、得られる厚鋼板のミクロ組織を適度に微細化するため、表面温度で900℃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、得られる厚鋼板において所望の靭性、表層硬さを確保できなくなる。このため、表面温度で900℃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33%以上である。
Cumulative rolling reduction in the temperature range of 900 ° C or less at the surface temperature: 30% or more In the present invention, in order to appropriately refine the microstructure of the resulting thick steel plate, the controlled rolling is performed in the temperature range of 900 ° C or less at the surface temperature. Do. If the cumulative reduction in the temperature range is less than 30%, the structure becomes coarse and the hardenability increases too much, so that the desired toughness and surface hardness cannot be secured in the resulting thick steel plate. For this reason, the cumulative reduction amount in the temperature range of 900 ° C. or less at the surface temperature is limited to 30% or more. In addition, Preferably it is 33% or more.

圧延終了温度:表面温度で870℃以下Ar3変態点以上
圧延終了温度が表面温度で870℃を超えると,組織が粗大化し、焼入性が増加しすぎて、得られる厚鋼板において所望の靭性、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870℃以下Ar3温度以上に限定した。なお、好ましくは780〜850℃である。
Rolling end temperature: 870 ° C or less at the surface temperature and above Ar3 transformation point If the rolling end temperature exceeds 870 ° C at the surface temperature, the structure becomes coarse and hardenability increases too much, and the desired toughness in the resulting thick steel plate, The surface hardness cannot be secured. On the other hand, if the rolling end temperature is less than the Ar3 transformation point at the surface temperature, ferrite is generated during 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 the surface temperature of 870 ° C or less and Ar3 temperature or more. In addition, Preferably it is 780-850 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。
Ar3変態点(℃)=900−332C+6Si−77Mn−20Cu−50Ni−18Cr−68Mo
(ここで、C、Si、Mn、Cu、Ni、Cr、Mo:各元素の含有量(質量%))
なお、上記式で記載された元素が含有されない場合には、当該元素を零として計算するものとする。
As the Ar3 transformation point, a value calculated using the following formula is used.
Ar3 transformation point (℃) = 900−332C + 6Si−77Mn−20Cu−50Ni−18Cr−68Mo
(Here, C, Si, Mn, Cu, Ni, Cr, Mo: content of each element (mass%))
In addition, when the element described by the said formula is not contained, the said element shall be calculated as zero.

圧延工程に引続いて、厚鋼板には、冷却工程が施される。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1/4t位置の平均冷却速度で3〜30℃/sの冷却速度で冷却し、表面温度が(Ar3変態点−100℃)以下400℃以上で、加速冷却を停止する冷却とする。
Subsequent to the rolling process, the steel plate is subjected to a cooling process. 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 first stage cooling starts from the temperature above the Ar3 transformation point at the surface temperature, cools at a cooling rate of 3 to 30 ° C / s at the average cooling rate at the thickness of 1/4 t, and the surface temperature is (Ar3 Transformation point −100 ° C.) Cooling that stops accelerated cooling at 400 ° C. or higher.

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り、表層部の靭性が低下する。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。
第一段冷却の冷却速度:板厚の1/4t位置の平均冷却速度で3〜30℃/s
冷却速度が3℃/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り、硬質相の生成量が不足して、所望のフェライトと硬質相との複相組織を得ることができず、所望の高強度を確保できなくなる。一方、冷却速度が30℃/sを超えて大きくなると、硬質相の生成量が増加し、降伏比が増加して、所望の低降伏比を確保できなくなる。このため、第一段冷却の冷却速度を、板厚の1/4t位置の平均冷却速度で3〜30℃/s冷却速度に限定した。なお、ここでいう「板厚の1/4t位置の平均冷却速度」とは、板厚1/4t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。
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 forms before the accelerated cooling starts and becomes coarse. Therefore, the fineness of the surface layer cannot be achieved, and the toughness of the surface layer portion decreases. For this reason, the start temperature of the first stage cooling was limited to the Ar3 transformation point or higher.
Cooling rate of the first stage cooling: 3-30 ° C / s at the average cooling rate at 1 / 4t position of the plate thickness
If the cooling rate is less than 3 ° C / s, the cooling is slow, and coarse and low toughness ferrite grains may be generated during cooling, the amount of hard phase generated is insufficient, and the desired phase of ferrite and hard phase is doubled. A tissue cannot be obtained, and a desired high strength cannot be ensured. On the other hand, when the cooling rate increases beyond 30 ° C./s, the amount of hard phase generated increases, the yield ratio increases, and the desired low yield ratio cannot be ensured. For this reason, the cooling rate of the first stage cooling was limited to a cooling rate of 3 to 30 ° C./s as an average cooling rate at a 1/4 t position of the plate thickness. Here, the “average cooling rate at the 1/4 t position of the plate thickness” means an average cooling rate from the start to the end of the accelerated cooling at the 1/4 t position of the plate thickness.

第一段冷却の冷却停止温度:表面温度で(Ar3変態点−100℃)以下400℃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で、表層部にフェライトを生成させる。これにより、表層部の硬さを低減でき、板厚中央部との硬さの差を小さくできる。冷却を停止する温度が、表面温度で(Ar3変態点−100℃)を超える温度では、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、冷却を停止する温度が400℃未満では、表層部の温度が低温となりすぎて、冷却中に相変態がほぼ完了してしまい、表層部はベイナイトやマルテンサイトなどの硬質相主体となる。
Cooling stop temperature for first stage cooling: (Ar3 transformation point –100 ° C) or less at surface temperature 400 ° C or higher In the first stage cooling in the present invention, cooling is performed so that the temperature difference between the surface layer part and the inside is larger. Then, ferrite is generated in the surface layer portion by recuperation after stopping the first stage cooling. Thereby, the hardness of a surface layer part can be reduced and the difference in hardness with a plate | board thickness center part can be made small. If the temperature at which the cooling is stopped 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, when the temperature at which the cooling is stopped is less than 400 ° C., the temperature of the surface layer portion becomes too low and the phase transformation is almost completed during cooling, and the surface layer portion is mainly composed of a hard phase such as bainite or martensite.

このようなことから、第一段冷却の冷却停止温度を、表面温度で(Ar3変態点−100℃)以下400℃以上の範囲に限定した。なお、好ましくは650〜450℃である。
また、本発明では、第一段冷却を、上記した1回の加速冷却からなる冷却に代えて、加速冷却を、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してもよい。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。これにより、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。
For this reason, the cooling stop temperature of the first stage cooling was limited to the surface temperature (Ar3 transformation point−100 ° C.) or lower and 400 ° C. or higher. In addition, Preferably it is 650-450 degreeC.
Further, in the present invention, the first-stage cooling may be the cooling that is repeated a plurality of times with the cooling stop and the subsequent recuperation sandwiched therebetween, instead of the cooling that consists of the single accelerated cooling described above. 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. Thereby, the choice of cooling temperature control can be expanded and the precision of cooling temperature control can be improved.

加速冷却を、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1/4t位置の平均冷却速度で2℃/s以上の冷却速度で、冷却停止温度が表面温度で400℃以上となる冷却を、冷却停止とその後の復熱とを挟み、複数回繰り返す冷却とすることが好ましい。なお、この冷却では、冷却停止温度が表面温度で(Ar3変態点−100℃)以下400℃以上となる加速冷却を少なくとも1回含むこととする。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図8に示す。   Accelerated cooling is repeated multiple times with cooling stop and subsequent recuperation, cooling starts from the surface temperature above the Ar3 transformation point, and the average cooling rate at the thickness of 1 / 4t It is preferable that the cooling at a cooling rate of 2 ° C./s or more and the cooling stop temperature at a surface temperature of 400 ° C. or higher is repeated a plurality of times with the cooling stop and the subsequent recuperation interposed therebetween. In this cooling, it is assumed that the cooling stop temperature includes accelerated cooling at a surface temperature of (Ar3 transformation point−100 ° C.) or lower and 400 ° C. or higher at least once. An example of the history of the steel sheet temperature when such cooling is performed is schematically shown in FIG.

第一段冷却における最初の加速冷却(第1回冷却)の開始温度:表面温度でAr3変態点以上
第一段冷却を構成する1回以上の加速冷却のうち、第1回冷却の開始温度は、表面温度でAr3変態点以上とすることが好ましい。第1回冷却の開始温度が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、表層部の靭性が低下する。このため、最初の加速冷却(第1回冷却)の開始温度を表面温度でAr3変態点以上に限定することが好ましい。
Start temperature of the first accelerated cooling (first cooling) in the first stage cooling: Ar3 transformation point or more at the surface temperature Among the one or more accelerated coolings constituting the first stage cooling, the starting temperature of the first cooling is The surface temperature is preferably not less than the Ar3 transformation point. If the first cooling start temperature is less than the Ar3 transformation point, ferrite is generated and coarsened before the start of accelerated cooling, so that the toughness of the surface layer portion decreases. For this reason, it is preferable to limit the starting temperature of the first accelerated cooling (first cooling) to the Ar3 transformation point or higher at the surface temperature.

第一段冷却における加速冷却の冷却速度:板厚1/4t位置の平均冷却速度で3℃/s以上
冷却速度が3℃/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1/4t位置の平均冷却速度で3℃/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、表層の過冷却を防ぐために30℃/s以下とすることが好ましい。鋼板の冷却速度は、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。ここでいう「板厚1/4t位置の平均冷却速度」とは、板厚1/4t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。図8で示せば、A点からB点までの平均の冷却速度をいう。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。
Cooling rate of accelerated cooling in the first stage cooling: 3 ° C / s or more as the average cooling rate at the position of 1 / 4t thickness When the cooling rate is less than 3 ° C / s, the cooling is slow and the ferrite grains are coarse and low toughness during cooling May generate. For this reason, the cooling rate of the first stage cooling was limited to 3 ° C./s or more in terms of the average cooling rate at the position of 1/4 t thickness. The upper limit of the cooling rate of the first stage cooling is not particularly limited, and is preferably 30 ° C./s or less in order to prevent supercooling of the surface layer. The cooling rate of the steel plate is substantially 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. Here, the “average cooling rate at the thickness 1/4 t position” refers to the average cooling rate from the start to the end of the accelerated cooling at the thickness 1/4 t position. If it shows in FIG. 8, it will say the average cooling rate from A point to B point. 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加速冷却の冷却停止温度:表面温度で400℃以上
第一段冷却における複数回の加速冷却において、表層部が400℃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。このため、すべての加速冷却の冷却停止温度を400℃以上に限定した。
冷却停止温度が表面温度で(Ar3変態点−100℃)以下400℃以上となる加速冷却:少なくとも1回
第一段冷却では、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、冷却停止後の復熱により、表層部にフェライトを生成させる。第一段冷却での複数回の加速冷却すべてにおいて、冷却停止温度が表面温度で(Ar3変態点−100℃)を超える温度では、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回を、冷却停止温度が(Ar3変態点−100℃)以下となる加速冷却とした。
Cooling stop temperature of accelerated cooling in the first stage cooling: 400 ° C or more at the surface temperature In multiple times of accelerated cooling in the first stage cooling, if the surface layer becomes less than 400 ° C, bainite and martensite transformation occurs during cooling As a result, the surface layer is hardened. For this reason, the cooling stop temperature of all accelerated cooling was limited to 400 ° C. or higher.
Cooling stop temperature at the surface temperature (Ar3 transformation point –100 ° C) or less 400 ° C or more Accelerated cooling: at least 1 time In the first stage cooling, cooling is performed so that a temperature difference between the surface layer part and the inside occurs to some extent. Ferrite is generated in the surface layer by recuperation after stopping. In all of the multiple times of accelerated cooling in the first stage cooling, if the cooling stop temperature exceeds the surface temperature (Ar3 transformation point-100 ° C), the steel plate temperature becomes too high during the subsequent reheating, and the surface layer part Insufficient ferrite formation occurs. For this reason, at least one of the plurality of accelerated coolings was set to accelerated cooling at which the cooling stop temperature was (Ar3 transformation point−100 ° C.) or lower.

本発明では、上記した第一段冷却を停止したのち、厚鋼板の表面温度が所定の温度まで復熱させたのち、第二段冷却を開始する。復熱は、表面温度で(Ar3変態点+10℃)以下650℃以上、表面と板厚中央の温度差が80℃以下となる時点まで行う。そして、第二段冷却を開始する。
本発明では、復熱中あるいは復熱後に、とくに第一段冷却と第二段冷却の間の冷却途中停止中の復熱後に、フェライトを生成させるため、復熱後の鋼板温度、すなわち第二段冷却の冷却開始温度が、フェライト生成量という組織制御の観点から重要な因子となる。
In the present invention, after the first stage cooling is stopped, the surface temperature of the thick steel plate is reheated to a predetermined temperature, and then the second stage cooling is started. Reheating is carried out until the surface temperature (Ar3 transformation point + 10 ° C) is 650 ° C or higher and the temperature difference between the surface and the center of the plate thickness is 80 ° C or lower. Then, the second stage cooling is started.
In the present invention, in order to generate ferrite during recuperation or after recuperation, particularly after recuperation during the cooling stop between the first stage cooling and the second stage cooling, the steel plate temperature after recuperation, that is, the second stage The cooling start temperature of the cooling is an important factor from the viewpoint of the structure control of the ferrite generation amount.

復熱が、表面温度で650℃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してしまい、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。また、復熱が、表面温度で(Ar3変態点+10℃)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り、表層部の延性が低下して部材変形能が低下する。   If the recuperation is less than 650 ° C. at the surface temperature, acicular ferrite and bainite with relatively high strength and yield ratio are generated in the surface layer, resulting in a decrease in the elongation of the surface layer and an increase in the yield ratio. The performance is reduced. Also, if the recuperation exceeds the surface temperature (Ar3 transformation point + 10 ° C), the phase transformation does not proceed after recuperation, ferrite formation in the surface layer becomes insufficient, the ductility of the surface layer decreases, and the member deforms The performance drops.

また、表面と板厚中央の温度差は、板厚方向のフェライト生成量の差を生じる原因となる。第二段冷却開始時点での、表面と板厚中央の温度差が80℃を超えると、表層部と板厚中央部とのミクロ組織差が大きくなりすぎ、大きな硬度差を生じる場合がある。板厚方向の硬度差が大きすぎると、地震などの変形時に部材としての変形能が低下する。
このようなことから、復熱後の温度、すなわち第二段冷却の冷却開始温度を、表面温度で(Ar3変態点+10℃)以下650℃以上、表面と板厚中央の温度差が80℃以下の範囲に限定した。
Further, the temperature difference between the surface and the center of the plate thickness causes a difference in the amount of ferrite generated in the plate thickness direction. If the temperature difference between the surface and the plate thickness center at the start of the second stage cooling exceeds 80 ° C., the microstructure difference between the surface layer portion and the plate thickness center portion becomes too large, and a large hardness difference may occur. When the hardness difference in the plate thickness direction is too large, the deformability as a member is reduced during deformation such as an earthquake.
For this reason, the temperature after recuperation, that is, the cooling start temperature of the second stage cooling, is 650 ° C or more at the surface temperature (Ar3 transformation point + 10 ° C), and the temperature difference between the surface and the center of the plate thickness is 80 ° C or less. It was limited to the range.

本発明の第二段冷却では、表面温度が(Ar3変態点+10℃)以下650℃以上、表面と板厚中央の温度差が80℃以下と、復熱したのち冷却を開始する。第二段冷却では、板厚の1/4t位置の平均冷却速度で3℃/s以上の冷却速度で、該第二段冷却を停止した後の復熱で表面温度が600℃以下になる冷却停止温度まで加速冷却する。
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+(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。
In the second stage cooling according to the present invention, the surface temperature is (Ar3 transformation point + 10 ° C.) or lower and 650 ° C. or higher, and the temperature difference between the surface and the center of the plate thickness is 80 ° C. or lower. In the second stage cooling, the cooling at which the surface temperature becomes 600 ° C. or less by the recuperation after stopping the second stage cooling at the cooling rate of 3 ° C./s or more at the average cooling rate at the 1/4 t position of the plate thickness. Accelerate cooling to stop temperature.
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 + (bainite and / or martensite)), and a desired high strength and low yield ratio can be realized.

第二段冷却の冷却速度:板厚の1/4t位置の平均冷却速度で3℃/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、3℃/s以上、好ましくは5℃/s以上で冷却する。冷却速度が3℃/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。
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℃超えとなるような温度では、硬質相への変態量が低下したり、自己焼戻しにより強度が低下し、所望の高強度を確保できなくなる。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℃以下になる温度に限定した。なお、好ましくは冷却停止温度は300℃以上とすることが所望の伸びと靭性を確保する観点から好ましい。復熱後の温度は、加速冷却停止時の板厚1/2t位置の温度に依存するので、各種伝熱計算による板厚1/2t位置の冷却停止温度から予測することができる。
Cooling rate of second stage cooling: 3 ° C / s or more at an average cooling rate at 1/4 ton of the plate thickness In order to make the untransformed part a hard phase, 3 ° C / s or more, preferably in second stage cooling Cool at 5 ° C / s or higher. If the cooling rate is less than 3 ° C./s, the amount of transformation to the hard phase decreases, and the desired high strength and low yield ratio cannot be realized.
Cooling stop temperature of the second stage cooling: Temperature at which the surface temperature becomes 600 ° C or less due to recuperation after stopping cooling Cooling stop temperature of the second stage cooling is the surface due to recuperation after stopping cooling of the second stage cooling If the temperature exceeds 600 ° C., 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. 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 less by recuperation after stopping the cooling. The cooling stop temperature is preferably 300 ° C. or more from the viewpoint of securing desired elongation and toughness. Since the temperature after recuperation depends on the temperature at the plate thickness 1 / 2t position when the accelerated cooling is stopped, it can be predicted from the cooling stop temperature at the plate thickness 1 / 2t position by various heat transfer calculations.

また、本発明では、第一段冷却と同様に、第二段冷却を、上記した1回の加速冷却からなる冷却に代えて、加速冷却を、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してもよい。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。
第二段冷却を構成する複数回の加速冷却のうち、第1回冷却は、表面温度が(Ar3変態点+10℃)(より好ましくはAr3変態点)以下650℃以上、表面と板厚中央の温度差が80℃以下に復熱したのち、冷却を開始することが好ましい。
Further, in the present invention, similarly to the first stage cooling, the second stage cooling is replaced with the cooling consisting of the above-described one-time accelerated cooling, and the accelerated cooling is sandwiched between the cooling stop and the subsequent recuperation, It is good also as cooling repeated several times. 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.
Of the multiple accelerated coolings that make up the second stage cooling, the first cooling has a surface temperature of (Ar3 transformation point + 10 ° C) (more preferably Ar3 transformation point) or less, 650 ° C or more, and surface and center of plate thickness It is preferable to start cooling after the temperature difference is reheated to 80 ° C. or less.

第二段冷却における第1回冷却の開始温度が、表面温度で(Ar3変態点+10℃)を超えると、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。また、開始温度が、表面温度で650℃未満では、表層部において針状フェライトやベイナイトが生成してしまいフェライト生成が不十分となり、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。このため、第二段冷却における最初の加速冷却(第1回冷却)の開始温度を表面温度で(Ar3変態点+10℃)以下(より好ましくはAr3変態点以下)650℃以上に限定することが好ましい。   If the starting temperature of the first cooling in the second stage cooling exceeds the surface temperature (Ar3 transformation point + 10 ° C.), the phase transformation does not proceed, and ferrite formation in the surface layer becomes insufficient. Also, if the starting temperature is less than 650 ° C. at the surface temperature, acicular ferrite and bainite are generated in the surface layer portion, resulting in insufficient ferrite generation, resulting in a decrease in the elongation of the surface layer portion and an increase in yield ratio. The performance is reduced. For this reason, it is possible to limit the starting temperature of the first accelerated cooling (first cooling) in the second stage cooling to 650 ° C. or more (more preferably below the Ar 3 transformation point) at the surface temperature (Ar 3 transformation point + 10 ° C.) or less. preferable.

また、表面と板厚中央の温度差は、板厚方向のフェライト生成量の差を生じる原因となる。第二段冷却における第1回冷却の開始時点での、表面と板厚中央の温度差が80℃を超えると、表層部と板厚中央部とのミクロ組織さが大きくなりすぎ、大きな硬度差を生じる場合がある。板厚方向の硬度差が大きすぎると、地震などの変形時に部材としての変形性能が低下する。   Further, the temperature difference between the surface and the center of the plate thickness causes a difference in the amount of ferrite generated in the plate thickness direction. If the temperature difference between the surface and the center of the plate thickness exceeds 80 ° C at the start of the first cooling in the second stage cooling, the microstructure of the surface layer portion and the center portion of the plate thickness becomes too large, resulting in a large hardness difference. May occur. When the hardness difference in the plate thickness direction is too large, the deformation performance as a member is deteriorated during deformation such as an earthquake.

このようなことから、第二段冷却の第1回冷却の開始は、表面温度で(Ar3変態点+10℃)(より好ましくはAr3変態点)以下650℃以上、表面と板厚中央の温度差が80℃以下となった時点とすることが好ましい。
第二段冷却を構成する複数回の加速冷却の冷却速度は、板厚1/4t位置の平均冷却速度で3℃/s以上とすることが好ましい。
For this reason, the start of the first cooling of the second stage cooling is the surface temperature (Ar3 transformation point + 10 ° C) (more preferably Ar3 transformation point) or less, 650 ° C or more, temperature difference between the surface and the center of the plate thickness It is preferable to set the temperature at 80 ° C. or lower.
The cooling rate of the plurality of times of accelerated cooling constituting the second stage cooling is preferably 3 ° C./s or more in terms of the average cooling rate at the position where the thickness is 1/4 t.

板厚1/4t位置の平均冷却速度で、冷却速度が3℃/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り、ベイナイト、マルテンサイト等の硬質相への変態量が低下し、所望の高強度と低降伏比を確保できなくなる。このため、第二段冷却の冷却速度を、板厚1/4t位置の平均冷却速度で3℃/s以上に限定した。なお、好ましくは6℃/s以上である。第二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る。なお、ここでいう「板厚1/4t位置の平均冷却速度」とは、板厚1/4t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。図8で示せば、C点からD点までの平均の冷却速度をいう。C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、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。   If the cooling rate is less than 3 ° C / s at an average cooling rate at a thickness of 1/4 t, the cooling is slow, and coarse and low toughness ferrite grains may be generated during cooling, and hard phases such as bainite and martensite. The amount of transformation to lowers and the desired high strength and low yield ratio cannot be ensured. For this reason, the cooling rate of the second stage cooling was limited to 3 ° C./s or more in terms of the average cooling rate at the position where the thickness is 1/4 t. In addition, Preferably it is 6 degrees C / s or more. The upper limit of the cooling rate of the second stage cooling is not particularly limited, and is almost determined by the plate thickness and the capacity of the cooling device. Here, the “average cooling rate at the position where the thickness is 1/4 t” refers to the average cooling rate from the start to the end of the accelerated cooling at the position where the thickness is 1/4 t. If it shows in FIG. 8, it will say the average cooling rate from C point to D point. Point C is the time when the temperature at the plate thickness 1 / 4t position 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.

加速冷却を停止し、復熱させたのちに、さらに上記した平均冷却速度の加速冷却を複数回繰り返して行う。そして、冷却停止後の復熱で表面温度が600℃以下になるような冷却停止温度まで加速冷却する冷却を、最終冷却として、第二段冷却を停止する。
第二段冷却の最終冷却が、冷却停止後の復熱で表面温度が600℃を超える冷却では、硬質相への変態量が少なく、さらに自己焼戻による強度低下により、所望の高強度を確保できなくなる。このようなことから、第二段冷却の最終冷却を、冷却停止後の復熱で表面温度が600℃以下になるような冷却停止温度まで加速冷却する冷却とすることが好ましい。
After the accelerated cooling is stopped and reheated, the above-described accelerated cooling at the average cooling rate is repeated a plurality of times. Then, the second-stage cooling is stopped by setting the cooling that is accelerated to the cooling stop temperature such that the surface temperature becomes 600 ° C. or less by the recuperation after stopping the cooling as the final cooling.
In the final cooling of the second stage cooling, when the surface temperature exceeds 600 ° C by recuperation after cooling stops, the transformation to the hard phase is small, and the desired high strength is ensured by the strength reduction due to self-tempering. become unable. For this reason, it is preferable that the final cooling of the second stage cooling is accelerated cooling to a cooling stop temperature at which the surface temperature becomes 600 ° C. or less by reheating after the cooling stop.

冷却停止温度の下限は特に限定する必要はないが、所望の伸びと靭性を確保する観点から、冷却停止後の復熱後で300℃以上にすることが好ましい。
なお、上記した冷却工程を施したのち、必要に応じて、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。
The lower limit of the cooling stop temperature is not particularly limited, but is preferably set to 300 ° C. or higher after the recuperation after the cooling stop from the viewpoint of securing desired elongation and toughness.
In addition, after performing the above-mentioned cooling process, you may perform a tempering process for the purpose of intensity | strength and toughness adjustment as needed. 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.

以下、実施例に基づいてさらに本発明について説明する。   Hereinafter, the present invention will be further described based on examples.

(実施例1)
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:50mm、あるいは25mm、80mmの厚鋼板とした。
なお、冷却工程は、第一段冷却と、冷却停止−復熱を経て、第二段冷却とからなる加速冷却とした。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、これに基づき、必要に応じて、板厚の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。
Example 1
The steel material having the composition shown in Table 1 was subjected to the rolling process and the cooling process shown in Table 2 to obtain thick steel sheets having a plate thickness of 50 mm, 25 mm, or 80 mm.
In addition, the cooling process was made into the accelerated cooling which consists of 1st stage cooling and 2nd stage cooling through cooling stop-recuperation. In each process, the steel plate temperature is measured by measuring the surface temperature with an infrared radiation thermometer, and based on this, various heat transfer calculation methods are used to calculate the temperature at the 1/4 t position of the plate thickness and the center thickness of the plate thickness. Used to calculate.

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
得られた厚鋼板から組織観察用試験片を採取し、L方向断面を研磨、ナイタール腐食して、表面から板厚方向に1〜5mmの領域である表層部と、板厚中央位置から±2mmの領域である板厚中央部について、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。また、表層部については、フェライトの公称粒径(平均結晶粒径)を求めた。フェライトの公称粒径(平均結晶粒径)は、各結晶粒の平均面積を求め、得られた平均面積の平方根をその厚鋼板のフェライト公称粒径(平均結晶粒径)とした。
(2)硬さ試験
得られた厚鋼板から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、鋼板表面から板厚方向に1〜5mmの領域(表層部)、および板厚中央位置から±2mmの領域(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10kg(98kN)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
得られた厚鋼板から、引張方向がL方向となるように、板厚の1/4t位置から、JIS Z 2201の規定に準拠して、JIS4号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
得られた厚鋼板の板厚1/4t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。
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 A specimen for microstructural observation is collected from the obtained thick steel plate, the cross section in the L direction is polished and subjected to nital corrosion, and the surface layer portion which is a region of 1 to 5 mm from the surface in the thickness direction, and the thickness Using the optical microscope (magnification: 400 times) or scanning electron microscope (magnification: 2000 times), the microstructure is observed over 3 fields of view and imaged at the center of the plate thickness, which is an area ± 2 mm from the center position. By image analysis, the type of structure and the structure fraction (area ratio) of ferrite were obtained. For the surface layer portion, the nominal grain size (average crystal grain size) of ferrite was determined. For the nominal grain size (average crystal grain size) of ferrite, the average area of each crystal grain was determined, and the square root of the obtained average area was defined as the nominal ferrite grain size (average crystal grain size) of the thick steel plate.
(2) Hardness test Take a specimen for hardness measurement from the obtained thick steel plate, and measure the hardness of the cross section in the thickness direction using a Vickers hardness tester in accordance with the provisions of JIS Z 2244. went. The measurement position is an area of 1 to 5 mm (surface layer part) in the sheet thickness direction from the steel sheet surface, and an area of ± 2 mm (sheet thickness center part) from the sheet thickness center position, with a 1 mm pitch in the sheet thickness direction in each area. Four or more points were measured. The test load (test force) was 10 kg (98 kN). The obtained hardness HV was arithmetically averaged to obtain the average hardness HV in that region.
(3) Tensile test JIS No. 4 tensile test specimens were collected from the obtained thick steel plate in accordance with the provisions of JIS Z 2201, from the 1/4 t position of the plate thickness, so that the tensile direction was the L direction. In accordance with JIS Z 2241, a tensile test was performed to determine tensile properties (yield strength YS, tensile strength TS). Further, the yield ratio YR (= YS / TS × 100%) was calculated from the obtained measured values.
(4) Impact test V-notch impact test specimens were collected in accordance with JIS Z 2242 from the position of 1 / 4t thickness and 1mm below the surface (the center position of the specimen was 6mm below the surface). Then, a Charpy impact test was carried out 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.

なお、板厚50mmの厚鋼板については、さらに、図3に示すような両面各1層(表面側入熱320kJ/cm、裏面側入熱150 kJ/cm)のサブマージアーク溶接により突き合せ継手を作製し、ノッチ位置を、表面側ボンド部として、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た。試験温度は0℃とし、試験片3本の吸収エネルギーの平均値を、その厚鋼板の大入熱溶接熱影響部の靭性として評価した。   In addition, for thick steel plates with a thickness of 50 mm, butt joints are further formed by submerged arc welding with one layer on each side (front surface heat input 320 kJ / cm, back surface heat input 150 kJ / cm) as shown in FIG. The V-notch impact test piece was sampled according to JIS Z 2242 using the notch position as the surface side bond part, and a Charpy impact test was performed. The test temperature was 0 ° C., and the average value of the absorbed energy of the three test pieces was evaluated as the toughness of the heat-affected zone of the large heat input welding of the thick steel plate.

また、得られた25mmおよび50mmの厚鋼板を用いて、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。なお、角形鋼管の断面寸法は、25mm厚×400×400 mm、50mm厚×600×600 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。得られた角形鋼管(プレスコラム)を用いて、コラム曲げ試験(三点曲げ試験)を実施した。試験方法はつぎのとおりとした。
(5)コラム曲げ試験
図4に示すように、得られた角形鋼管(プレスコラム)(長さ:3250mm)1a,1a各々に、SN490鋼板製通しダイアフラム(板厚40mm)2a,2aを炭酸ガス溶接で溶接し、ついで、2枚のダイアフラム2a,2a間に4面BOX柱3aを配して炭酸ガス溶接し、コラム曲げ試験体とした。なお、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5に、試験体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。
Further, square steel pipes (press columns) were produced by cold pressing using the obtained 25 mm and 50 mm thick steel plates. The cross-sectional dimensions of the square steel pipe were 25 mm thickness x 400 x 400 mm, 50 mm thickness x 600 x 600 mm, and the seam welding was a submerged arc welding with one layer on each side. A column bending test (three-point bending test) was performed using the obtained rectangular steel pipe (press column). The test method was as follows.
(5) Column bending test As shown in Fig. 4, the obtained square steel pipe (press column) (length: 3250mm) 1a, 1a is passed through SN490 steel plate diaphragm (plate thickness 40mm) 2a, 2a with carbon dioxide gas. Welding was performed by welding, and then a four-sided BOX column 3a was placed between the two diaphragms 2a and 2a, and carbon dioxide gas welding was performed to obtain a column bending specimen. The strength and rigidity of the 4-sided BOX column were made sufficiently higher than that of the press column to prevent plastic deformation other than the press column during the test. FIG. 5 shows an enlarged view of the vicinity of the welded portion of the press column 1a, diaphragm 2a, and 4-sided BOX column 3a in the test body.

得られた試験体の両端部を支持し、図6に示すように、試験体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7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。
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(累積塑性変形角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を求めた。なお、ηは次式より算出される。
η=Σθpl/θp
ここで、θp=(Pp/2)L/(3・E・I)+Pp/2/(G・Aw)
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(=205000MPa)、G:剪断剛性率(=79000MPa)、
Mp:コラムの全塑性モーメント
Aw:剪断面積(mm
なお、コラムの全塑性モーメントMpは、次式で定義される。
As shown in FIG. 6, a three-point repeated bending test (column bending test) was performed in which positive and negative loads were repeatedly applied to the center of the test body in the vertical direction 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. 7 was created.
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 of the specimen (cumulative plastic deformation angle Σθpl) is obtained and tested. The cumulative plastic deformation ratio η was determined as an index of the plastic deformability of the body. Η is calculated from the following equation.
η = Σθpl / θp
Here, θp = (Pp / 2) L 2 / (3 · E · I) + Pp / 2 / (G · Aw)
Pp: Total plastic load (N) = Mp / L,
L: Column cantilever length (distance from diaphragm to column end support point (= 3250mm)
E: Young's modulus (= 205000 MPa), G: shear rigidity (= 79000 MPa),
Mp: Total plastic moment of column Aw: Shear area (mm 2 )
The total plastic moment Mp of the column is defined by the following equation.

Mp=I・σ/{D/√2−(√2−1)R}、
ここで、I:コラムの断面二次モーメント、σ:鋼材の降伏強さ(MPa)、
また、コラムの断面二次モーメントIは、次式で定義される。
I=(t/12){4(D−2R)(2D+14R+2t−2DR−3Dt−12Rt)+3π(2R−t)(2D+6R+t−4DR−2Rt)}
ここで、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(mm)、R=r+t
また、剪断面積Awは、次式で、定義される。
Mp = I · σ y / {D / √2− (√2−1) R},
Where, I: Column moment of inertia, σ y : Yield strength of steel (MPa),
The column cross-sectional secondary moment I is defined by the following equation.
I = (t / 12) { 4 (D-2R) (2D 2 + 14R 2 + 2t 2 -2DR-3Dt-12Rt) + 3π (2Rt) (2D 2 + 6R 2 + t 2 -4DR-2Rt)}
Here, D: column diameter (mm), t: column plate thickness (mm), r: bending radius (mm) of the inner surface of the column corner, R = r + t
The shear area Aw is defined by the following equation.

Aw=I/[r{((D−t−2r)/√2)・π/4+r/√2}
+(√2)/2{((D−t−r)/√2)−(r/√2)
+r(1−1/√2)]
なお、累積塑性変形倍率ηが30以上である場合、構造部材の耐震性(塑性変形能)に優れるとする。
Aw = I / [r {(((D−t−2r) / √2) · π / 4 + r / √2}
+ (√2) / 2 {((D−t−r) / √2) 2 − (r / √2) 2 }
+ R 2 (1-1 / √2)]
When the cumulative plastic deformation magnification η is 30 or more, it is assumed that the structural member is excellent in earthquake resistance (plastic deformability).

本発明例はいずれも、降伏強さYS:385MPa以上,引張強さTS:550MPa以上,降伏比YR:80%以下を有し、さらに表面層および板厚1/4t位置での破面遷移温度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となっている。さらに、本発明例はいずれも、表面側入熱:320kJ/cmのサブマージアーク溶接の溶接ボンド部の0℃における吸収エネルギーが70J以上と、大入熱溶接部靭性にも優れている。   Each of the inventive examples has a yield strength YS: 385 MPa or more, a tensile strength TS: 550 MPa or more, a yield ratio YR: 80% or less, and a fracture surface transition temperature at a surface layer and a thickness of 1/4 t. It is a high-strength, high-toughness, non-tempered, low yield ratio, high-tensile steel plate with vTrs of -40 ° C or lower. Furthermore, in all of the examples of the present invention, the absorbed energy at 0 ° C. of the weld bond portion of the submerged arc welding of surface side heat input: 320 kJ / cm is 70 J or more, and the high heat input weld zone toughness is excellent.

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する。また、本発明例はいずれも、冷間曲げ加工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(コラム曲げ試験)における累積塑性変形倍率ηが30以上であり、耐震性(塑性変形性能)に優れた、構造部材とすることができる、非調質低降伏比高張力厚鋼板であるといえる。一方、本発明の範囲を外れる比較例は、強度、降伏比、靭性が不足しているか、あるいは大入熱溶接部靭性が低下しているか、または、冷間加工後の表層部の延性、靭性が低下し、構造部材としての耐震性(塑性変形性能)が低下している。
(実施例2)
表1に示す鋼No.A,No.Eの組成を有する鋼素材に、表4に示す圧延工程、冷却工程を施し、板厚:50mmの厚鋼板とした。
Furthermore, all of the examples of the present invention have a thickness direction hardness distribution in which the average hardness of the surface layer portion 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. Also, in all of the examples of the present invention, when a cold column is processed and processed into a press column to constitute a press column-diaphragm joint structural member, a press column-diaphragm joint three-point bending test (column bending test) It can be said that it is a non-tempered low yield ratio high-tensile steel plate that has a cumulative plastic deformation ratio η of 30 or more and can be a structural member that is excellent in earthquake resistance (plastic deformation performance). On the other hand, the comparative examples out of the scope of the present invention are insufficient in strength, yield ratio, toughness, or low in high heat input weld toughness, or ductility and toughness of the surface layer after cold working The seismic resistance (plastic deformation performance) as a structural member is reduced.
(Example 2)
Steel No. A and No. 1 shown in Table 1. The steel material having the composition E was subjected to the rolling process and the cooling process shown in Table 4 to obtain a thick steel plate having a thickness of 50 mm.

なお、冷却工程は、第一段冷却と、冷却停止−復熱を経て、第二段冷却とからなる加速冷却とした。そして、第一段冷却を、間に冷却停止とその後の復熱とを含んだ1回以上の加速冷却からなる冷却とし、第二段冷却を間に冷却停止とその後の復熱とを含んだ1回以上の加速冷却からなる冷却とした。なお、実施例1と同様に、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、これに基づき、必要に応じて、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。   In addition, the cooling process was made into the accelerated cooling which consists of 1st stage cooling and 2nd stage cooling through cooling stop-recuperation. Then, the first stage cooling is a cooling consisting of one or more accelerated coolings including the cooling stop and the subsequent recuperation, and the second stage cooling includes the cooling stop and the subsequent recuperation. The cooling consisted of one or more accelerated coolings. In addition, like Example 1, the steel plate temperature in each process measures the surface temperature with an infrared radiation thermometer, and based on this, the temperature at the plate thickness 1 / 4t position and the plate thickness central temperature are set as necessary. It was calculated using various heat transfer calculation methods.

得られた厚鋼板について、実施例1と同様に、組織観察、硬さ試験、引張試験、衝撃試験を実施した。
得られた結果を表5に示す。
また、実施例1と同様に、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。なお、角形鋼管(プレスコラム)の断面寸法は、600×600mm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。
The obtained thick steel plate was subjected to structure observation, hardness test, tensile test, and impact test in the same manner as in Example 1.
The results obtained are shown in Table 5.
Further, in the same manner as in Example 1, a square steel pipe (press column) was produced by cold pressing using the obtained thick steel plate. In addition, the cross-sectional dimension of the square steel pipe (press column) was 600 × 600 mm, and the seam welding was a submerged arc welding with one layer on each side.

そして、実施例1と同様に、得られた角形鋼管(プレスコラム)(長さ3250mm)1a、1a各々に、SN490鋼板製通しダイアフラム(板厚40mm)2aを炭酸ガス溶接で溶接し、ついで、2枚のダイアフラム2a、2a間に4面BOX柱3aを配して炭酸ガス溶接し、コラム曲げ試験の試験体とした。
そして、実施例1と同様に、得られた試験体の両端部を支持し、図6に示すように、試験体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。そして、実施例1と同様に、荷重Pと変形量(回転角)θを測定し、図7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。
In the same manner as in Example 1, the obtained square steel pipe (press column) (length 3250 mm) 1a and 1a was welded with a SN490 steel plate through diaphragm (plate thickness 40 mm) 2a by carbon dioxide welding, A four-sided BOX column 3a was placed between the two diaphragms 2a and 2a, and carbon dioxide was welded to obtain a specimen for a column bending test.
And like Example 1, the both ends of the obtained test body are supported, and as shown in FIG. 6, the positive and negative load is repeatedly applied to the center part of the test body in the vertical direction (3-point repeated bending test ( Column bending test). Then, similarly to Example 1, 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. 7 was created.

そして、実施例1と同様に、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を求めた。
得られた結果を表5に併記した。
As in Example 1, when the load (moment) decreased by 5% from the maximum value due to local buckling or brittle fracture, the specimen was considered to be fractured, and the total plastic rotation angle of the specimen (accumulated) The plastic rotation angle Σθpl) was determined, and the cumulative plastic deformation ratio η was determined as an index of the plastic deformability of the specimen.
The obtained results are also shown in Table 5.

本発明例はいずれも、降伏強さYS:385MPa以上、引張強さTS:550MPa以上、降伏比YR:80%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となっている。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、非調質低降伏比高張力厚鋼板であるといえる。一方、本発明の範囲を外れる比較例は、所望の降伏比、所望の板厚方向硬さ分布が確保できていないか、構造部材としての累積塑性変形倍率が低くなっている。   Each of the inventive examples has a yield strength YS: 385 MPa or more, a tensile strength TS: 550 MPa or more, a yield ratio YR: 80% or less, and vTrs at the surface layer portion and 1/4 t position in the plate thickness direction is − It is a high-strength, high-toughness, non-tempered, low yield ratio, high-tensile steel plate that satisfies 40 ° C or lower. Furthermore, each of the inventive examples 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 center portion is 60 HV or less, and cold bending is performed. When a press column-diaphragm joint structural member is constructed 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 said that it is an excellent non-tempered low yield ratio high-tensile thick steel plate that can be used as a structural member. On the other hand, in a comparative example outside the scope of the present invention, a desired yield ratio and a desired thickness direction hardness distribution are not ensured, or the cumulative plastic deformation ratio as a structural member is low.

1 柱
1a プレスコラム
2 ダイアフラム
2a 通しダイアフラム
3 溶接部
3a 4面BOX柱
4 当金
4a コラム/ダイアフラム溶接部
5a ダイアフラム/4面BOX柱溶接部
1 Column 1a Press column 2 Diaphragm 2a Through diaphragm 3 Welded portion 3a 4 side BOX column 4 Metal 4a Column / diaphragm welded portion 5a Diaphragm / 4 side BOX column welded portion

Claims (9)

質量%で、
C:0.05〜0.10%、 Si:0.01〜0.45%、
Mn:1.2〜1.8%、 P:0.020%以下、
S:0.0010〜0.0030%、 Al:0.05%以下、
Ti:0.005〜0.020%、 N:0.0030〜0.0060%
を含み、TiとNを、Ti含有量とN含有量との比、Ti/Nが2.0〜4.0を満足するように含有し、さらに、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種以上を含有し、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%以下に制限し、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組成と、
少なくとも、鋼板表面から板厚方向に1〜5mmの表層部がフェライトと、硬質相としてパーライト、ベイナイト、マルテンサイトのうち1種または2種以上からなり、前記フェライトの平均結晶粒径が4.0〜18.0μmである組織を有し、
鋼板表面から板厚方向に1mm〜5mmの表層部の平均硬さが225HV以下で、該表層部と板厚中央位置を中心に±2mmの範囲である板厚中央部との硬度差が60HV以下である板厚方向硬さ分布を有し、冷間加工後の表層部延性・靭性に優れ、かつ大入熱溶接部靭性に優れることを特徴とする
降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:80%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:各元素の含有量(質量%)
% By mass
C: 0.05 to 0.10%, Si: 0.01 to 0.45%,
Mn: 1.2 to 1.8%, P: 0.020% or less,
S: 0.0010 to 0.0030%, Al: 0.05% or less,
Ti: 0.005-0.020%, N: 0.0030-0.0060%
Ti and N are contained so that the ratio of Ti content to N content, Ti / N satisfies 2.0 to 4.0, and 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: One or more selected from 0.0003 to 0.0030%, Nb and Mo as impurity elements, Nb: 0.004% Hereinafter, Mo is limited to 0.04% or less, the carbon equivalent Ceq defined by the following formula (1) satisfies 0.35 to 0.48, the composition composed of the balance Fe and inevitable impurities,
At least the surface layer portion of 1 to 5 mm in the thickness direction from the steel sheet surface is composed of ferrite and one or more of pearlite, bainite and martensite as the hard phase, and the average crystal grain size of the ferrite is 4.0 to 18.0. having a tissue that is μm,
The average hardness of the surface layer part of 1mm to 5mm from the surface of the steel sheet is 225HV or less, and the difference in hardness between the surface layer part and the sheet thickness center part within ± 2mm around the sheet thickness center position is 60HV or less. Yield strength: 385 MPa or higher, tensile strength: characterized by having a thickness distribution in the thickness direction, excellent surface layer ductility and toughness after cold working, and high heat input weld toughness Non-tempered low yield ratio high tensile steel plate with 550MPa or more and yield ratio: 80% or less.
Record
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
Here, C, Mn, Cr, Mo, V, Cu, Ni: Content of each element (mass%)
前記組成に加えてさらに、質量%で、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種以上を含有する組成と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。   2. The composition according to claim 1, wherein the composition further contains one or more selected from Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050% by mass% in addition to the composition. Non-tempered low yield ratio high-tensile thick steel plate as described in 1. 前記Ca、REMを、下記(2)式で定義されるACRが0.2〜0.8を満足するように含有することを特徴とする請求項2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:各元素の含有量(質量%)
The non-tempered low yield ratio high-tensile thick steel plate according to claim 2, wherein the Ca and REM are contained so that the ACR defined by the following formula (2) satisfies 0.2 to 0.8.
Record
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
Here, Ca, REM, O, S: Content of each element (mass%)
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、該厚鋼板に途中冷却停止を含む第一段冷却と第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、を施す非調質厚鋼板の製造方法において、
前記鋼素材を、質量%で、
C:0.05〜0.10%, Si:0.01〜0.45%、
Mn:1.2〜1.8%、 P:0.020%以下、
S:0.0010〜0.0030%、 Al:0.05%以下、
Ti:0.005〜0.020%、 N:0.0030〜0.0060%
を含み、TiとNを、Ti含有量とN含有量との比、Ti/Nが2.0〜4.0を満足するように含有し、さらに、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種以上を含有し、さらに、不純物元素としてNb、Moを、Nb:0.004%以下、Mo:0.04%以下に制限し、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組成を有する鋼素材とし、
前記熱間圧延の加熱温度を1050〜1200℃とし、
前記熱間圧延を、表面温度で900℃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870℃以下Ar3変態点以上となる圧延とし、
前記第一段冷却が、表面温度でAr3変態点以上の温度から冷却を開始し、板厚1/4t位置の平均冷却速度で3〜30℃/sの冷却速度で冷却し、表面温度が(Ar3変態点−100℃)以下400℃以上で、加速冷却を停止する冷却とし、
冷却停止後、複熱し、表面温度が(Ar3変態点+10℃)以下650℃以上、表面と板厚中央の温度差が80℃以下となる時点で、前記第二段冷却を開始し、
該第二段冷却を、板厚1/4t位置の平均冷却速度で3℃/s以上の冷却速度で、該第二段冷却を停止した後の復熱で表面温度が600℃以下になるような冷却停止温度まで加速冷却する冷却と、
することを特徴とする降伏強さ:385MPa以上、引張強さ:550MPa以上、降伏比:80%以下を有し、冷間加工後の表層部延性・靭性に優れ、かつ大入熱溶接部靭性に優れた非調質低降伏比高張力厚鋼板の製造方法。

Ceq=C+Mn/6+(Cr+Mo+V)/5+(Cu+Ni)/15 ‥‥(1)
ここで、C、Mn、Cr、Mo、V、Cu、Ni:各元素の含有量(質量%)
A steel material is subjected to a hot rolling process to make a thick steel plate, and following the rolling step, the thick steel plate is subjected to two-stage accelerated cooling consisting of first-stage cooling and second-stage cooling including stopping cooling on the way. In the manufacturing method of the non-tempered thick steel sheet,
The steel material in mass%,
C: 0.05-0.10%, Si: 0.01-0.45%,
Mn: 1.2 to 1.8%, P: 0.020% or less,
S: 0.0010 to 0.0030%, Al: 0.05% or less,
Ti: 0.005-0.020%, N: 0.0030-0.0060%
Ti and N are contained so that the ratio of Ti content to N content, Ti / N satisfies 2.0 to 4.0, and 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: One or more selected from 0.0003 to 0.0030%, Nb and Mo as impurity elements, Nb: 0.004% Hereinafter, Mo is limited to 0.04% or less, and the carbon equivalent Ceq defined by the following formula (1) satisfies 0.35 to 0.48, and the steel material has a composition composed of the balance Fe and inevitable impurities,
The heating temperature of the hot rolling is 1050 to 1200 ° C,
The hot rolling is a rolling in which the cumulative reduction amount in the temperature range of 900 ° C. or less at the surface temperature is 30% or more, and the rolling end temperature is 870 ° C. or less in the surface temperature and the Ar 3 transformation point or more,
The first stage cooling starts from a temperature above the Ar 3 transformation point at the surface temperature, cools at a cooling rate of 3 to 30 ° C / s at an average cooling rate at a thickness of 1/4 t, and the surface temperature is (Ar 3 transformation point-100 ° C) Below 400 ° C or higher, accelerated cooling is stopped,
After stopping the cooling, the second stage cooling is started when the surface temperature becomes (Ar 3 transformation point + 10 ° C.) or less 650 ° C. or more and the temperature difference between the surface and the plate thickness becomes 80 ° C. or less.
The second stage cooling is performed at a cooling rate of 3 ° C./s or more at an average cooling rate at a thickness of 1/4 t, and the surface temperature is reduced to 600 ° C. or less by recuperation after stopping the second stage cooling. Cooling to accelerate cooling to the proper cooling stop temperature,
Yield strength: 385 MPa or more, tensile strength: 550 MPa or more, yield ratio: 80% or less, excellent surface layer ductility and toughness after cold working, and high heat input weld toughness A method for producing non-tempered, low yield ratio, high-tensile thick steel plates with excellent resistance.
Record
Ceq = C + Mn / 6 + (Cr + Mo + V) / 5 + (Cu + Ni) / 15 (1)
Here, C, Mn, Cr, Mo, V, Cu, Ni: Content of each element (mass%)
前記第一段冷却に代えて、第一段冷却を、表面温度でAr3変態点以上の温度から冷却を開始し、板厚1/4t位置の平均冷却速度で3℃/s以上の冷却速度で、冷却停止温度が表面温度で400℃以上となる加速冷却を、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し、前記複数回の加速冷却が、冷却停止温度が表面温度で(Ar3変態点−100℃)以下400℃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4に記載の非調質低降伏比高張力厚鋼板の製造方法。   Instead of the first stage cooling, the first stage cooling is started from the temperature above the Ar3 transformation point at the surface temperature, and the average cooling rate at a thickness of 1/4 t is 3 ° C / s or more. The accelerated cooling at which the cooling stop temperature is 400 ° C. or more at the surface temperature is the cooling that is repeated a plurality of times across the cooling stop and the subsequent recuperation, and the plurality of accelerated coolings are performed at the surface temperature. The method for producing a non-tempered low yield ratio high-tensile thick steel plate according to claim 4, comprising accelerated cooling (Ar3 transformation point-100 ° C) to 400 ° C or higher at least once. 前記第二段冷却に代えて、第二段冷却を、板厚1/4t位置の平均冷却速度で3℃/s以上の冷却速度で、冷却停止とその後の復熱とを挟んで、加速冷却を複数回繰り返す冷却とし、前記複数回の加速冷却のうち、冷却停止後の復熱で表面温度が600℃以下になるような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4または5に記載の非調質低降伏比高張力厚鋼板の製造方法。   Instead of the above-mentioned second stage cooling, the second stage cooling is accelerated cooling with an average cooling rate at a thickness of 1/4 t at a cooling rate of 3 ° C./s or more and a cooling stop and subsequent recuperation. The cooling is repeated a plurality of times, and among the plurality of accelerated coolings, the final cooling is accelerated cooling that cools to a cooling stop temperature such that the surface temperature becomes 600 ° C. or less by reheating after the cooling stop. The manufacturing method of the non-tempered low yield ratio high-tensile thick steel plate of Claim 4 or 5. 前記冷却工程に引続き、400℃以上700℃以下の温度で焼戻しを行う焼戻工程を施すことを特徴とする請求項4ないし6のいずれかに記載の非調質低降伏比高張力厚鋼板の製造方法。   The tempering step of tempering at a temperature of 400 ° C or higher and 700 ° C or lower is performed subsequent to the cooling step, wherein the tempered low yield ratio high-tensile thick steel plate according to any one of claims 4 to 6 is provided. Production method. 前記鋼素材の組成に加えて、さらに、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種以上を含有する組成とすることを特徴とする請求項4ないし7のいずれかに記載の非調質低降伏比高張力厚鋼板の製造方法。   5. In addition to the composition of the steel material, the composition further includes one or more selected from Ca: 0.0005 to 0.0050% and REM: 0.0010 to 0.0050%. The manufacturing method of the non-tempered low yield ratio high tension thick steel plate in any one of thru | or 7. 前記Ca、REMを、下記(2)式で定義されるACRが0.2〜0.8を満足するように含有することを特徴とする請求項8に記載の非調質低降伏比高張力厚鋼板の製造方法。

ACR=[(Ca+0.29×REM)−{0.18+130×(Ca+0.29×REM)}×O]/(1.25×S) ‥‥ (2)
ここで、Ca、REM、O、S:各元素の含有量(質量%)
The said Ca and REM are contained so that ACR defined by following (2) Formula may satisfy 0.2-0.8, The manufacture of the non-tempered low yield ratio high-tensile steel plate of Claim 8 characterized by the above-mentioned. Method.
Record
ACR = [(Ca + 0.29 × REM) − {0.18 + 130 × (Ca + 0.29 × REM)} × O] / (1.25 × S) (2)
Here, Ca, REM, O, S: Content of each element (mass%)
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