JP2004003008A - Hot-rolled non-thermally refined bar steel and manufacturing method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a hot-rolled non-thermally refined bar steel and a manufacturing method therefor. <P>SOLUTION: This bar steel has a ferrite single-phase structure in which fine precipitates having particle sizes of less than 10 nm are dispersed, and has a composition comprising, by mass%, one or more elements of 0.15% or less C, 0.5% or less Si, 2% or less Mn, 0.1% or less Al, 0.03-0.35% Ti, 0.05-0.8% Mo, 0.08% or less Nb, 0.15% or less V, and 1.5% or less W, one or more elements of 0.03-0.1% S, 0.2% or less Pb, 0.005% or less Ca, and 0.02% or less B, as needed, so as to satisfy 0.5≤(C/12)/ä(Ti/48)+(Mo/96)+(Nb/93)+(V/51)+(W/192)}≤1.5, (wherein C, Ti, Mo, Nb, V, and W express each added quantity (mass%) while the quantity is 0 when those are not added), and the balance Fe with unavoidable impurities. The manufacturing method comprises steps for heating the steel having the composition at 1,100°C or higher, hot-forging it, and then cooling it to 550-700°C at a rate of 0.5°C/sec or lower, or holding it at 550-700°C for 10 minutes or longer after air-cooling it. <P>COPYRIGHT: (C)2004,JPO

Description

但し、各元素は含有量（質量％）とする。
【００１２】
４．微細析出物がＴｉ、Ｍｏの炭化物であることを特徴とする１乃至３のいずれか一つに記載の引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼。
【００１３】
５．鋼組成として、更に質量％で、Ｎｂ≦０．０８％、Ｖ≦０．１５％、Ｗ≦１．５％の一種または二種以上を含有する２記載の引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼。
【００１４】
６．鋼組成として更に式（２）を満足することを特徴とする５記載の引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼。
０．５≦（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）
＋（Ｖ／５１）＋（Ｗ／１８４）｝≦１．５　　　　（２）
但し、各元素は含有量（％）とし、含まれないものは０とする。
【００１５】
７．微細析出物がＴｉとＭｏとＮｂ、Ｖ、Ｗの内の少なくとも一種とを含む炭化物であることを特徴とする５または６記載の引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼。
【００１６】
８．鋼組成として更に質量％で、Ｓ：０．０３〜０．１％、Ｐｂ≦０．２％、Ｃａ≦０．００５％、Ｂ≦０．０２％の一種または二種以上を含有することを特徴とする２、３、５、６のいずれか一つに記載の引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼。
【００１７】
９．鋼組成が２、３、５、６、８のいずれか一つに記載の鋼を１１００℃以上に加熱後、仕上げ圧延温度８００℃以上で熱間圧延し、その後の冷却において、７００〜５５０℃を０．５℃／ｓｅｃ以下の冷却速度で冷却することを特徴とする引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼の製造方法。
【００１８】
１０．鋼組成が２、３、５、６、８のいずれか一つに記載の鋼を１１００℃以上に加熱後、仕上げ圧延温度８００℃以上で熱間圧延し、その後の冷却において、７００〜５５０℃を０．５℃／ｓｅｃ超えの冷却速度で冷却し、その後５５０〜７００℃で１０分間以上再加熱することを特徴とする引張強さ７００ＭＰａ以上で０．８５以上の降伏比を有する熱間圧延型非調質棒鋼の製造方法。
【００１９】
【発明の実施の形態】
本発明のミクロ組織、成分組成および製造条件について以下に詳細に説明する。
【００２０】
１．ミクロ組織
本発明鋼はそのミクロ組織をフェライト単相組織に粒径１０ｎｍ未満の微細析出物を含む組織に規定する。
【００２１】
フェライト単相組織とした場合、調質材に匹敵する靭性が得られ、該組織中に微細析出物を分散析出させた場合、調質材に匹敵する高降伏強度が被削性を損なうことなく得られる。微細析出物は熱間圧延後の冷却速度の調整や、熱間圧延後の析出処理により析出させる。
【００２２】
本発明では微細析出物は粒径１０ｎｍ未満とする。析出物の粒径が１０ｎｍ以上の場合、輸送機器、建機の機械構造用部品として必要な引張強さ７００ＭＰａ以上が得られない。
【００２３】
また、フェライト単相組織中に粒径１０ｎｍ未満の微細析出物を析出させた場合、降伏比が０．８５以上となり、降伏強度の上昇に対して引張強度の上昇が抑えられ鋼の硬化が小さく調質鋼に匹敵する被削性が得られる。母相、析出物のいずれかが本発明の規定外となった場合、降伏比は０．８５未満となる。
【００２４】
微細析出物の粒径は小さいほど強度向上に有効で、望ましくは５ｎｍ，更に望ましくは３ｎｍ以下とし、そのような微細析出物としてＴｉ、Ｍｏを複合含有した炭化物、またそれらに更にＮｂ、Ｖ、Ｗの一種または二種以上を含む炭化物が好ましい。
【００２５】
これらの微細析出物の分布形態は特に規定しないが、母相中に均一分散（分散析出）することが望ましい。
【００２６】
また、本発明において、微細析出物の大きさは、全析出物の９０％以上で満足すれば、目的とする引張強さ７００ＭＰａ以上が得られる。但し、１０ｎｍ以上の大きさの析出物は析出物形成元素を消費し、強度に悪影響をあたえるため、５０ｎｍ以下とすることが好ましい。
【００２７】
上述した析出物とは別に少量のＦｅ炭化物を含有しても本発明の効果は損なわれないが、平均粒径が１μｍ以上のＦｅ炭化物を多量に含むと靭性を阻害するため、本発明においては含有されるＦｅ炭化物の大きさ上限は１μｍ、含有率は全体の１％以下とすることが望ましい。
【００２８】
微細析出物の全析出物に占める割合は、以下の方法により求める。電子顕微鏡試料を、ツインジェット法を用いた電解研磨法で作成し、加速電圧２００ｋＶで観察する。その際、微細析出物が母相に対して計測可能なコントラストになるように母相の結晶方位を制御し、析出物の数え落としを最低限にするために焦点を正焦点からずらしたデフォーカス法で観察を行う。
【００２９】
また、析出物粒子の計測を行った領域の試料の厚さは電子エネルギー損失分光法を用いて、弾性散乱ピークと非弾性散乱ピーク強度を測定することで評価する。
【００３０】
この方法により、粒子数の計測と試料厚さの計測を同じ領域について実行することができる。粒子数および粒子径の測定は試料の０．５×０．５μｍの領域４箇所について行い、１μｍ^２当たりに分布する析出物を粒径ごとの個数として算出する。
【００３１】
この値と試料厚さから、析出物の１μｍ^３当たりに分布する粒子径ごとの個数を算出し、径が１０ｎｍ未満の析出物について、測定した全析出物に占める割合を算出する。
【００３２】
また、本発明においてフェライト単相組織とは、断面組織観察（２００倍の光学顕微鏡組織観察）でフェライト面積率９５％以上とし、好ましくは９８％以上とする。
【００３３】
２．成分組成
本発明鋼は上述したミクロ組織で目的とする性能が得られるが、以下の成分組成が好ましい。
【００３４】
Ｃ
Ｃは０．１５％を超えて含有すると微細析出物が粗大化し、強度が低下するため０．１５％以下とすることが好ましい。より好ましくは０．０３％以上０．１２％以下である。
【００３５】
Ｓｉ
Ｓｉは冷間加工性を向上させるため添加する。０．５％をこえるとその効果が損なわれるようになるため、０．５％以下とする。より好ましくは０．１５％以下である。
【００３６】
Ｍｎ
Ｍｎは強度、延性を向上させるため添加する。２％を超えるとその効果が損なわれるため２％以下とする。より好ましくは０．５％以上１．８％以下である。
【００３７】
Ａｌ
Ａｌは脱酸剤として作用する。またＮとＡｌＮを形成し、Ｂの焼入れ性効果を向上させる。０．１％を超えるとその効果が飽和するため０．１％以下とする。より好ましくは０．０５％以下である。
【００３８】
Ｔｉ
ＴｉはＴｉ系炭化物、ＭｏとともにＴｉ−Ｍｏ系炭化物を含む析出物を微細に析出させ、強度を向上させるため添加する。引張強度７００ＭＰａ以上を確保するため０．０３％以上とし、一方、０．３５％を超えて添加すると析出物が粗大化し、強度、靭性が低下するため０．０３〜０．３５％とする。より好ましくは０．０３〜０．２０％である。
【００３９】
Ｍｏ
ＭｏはＭｏ系炭化物、ＴｉとともにＴｉ−Ｍｏ系炭化物を含む析出物を微細に析出させ、強度を向上させるため添加する。引張強度７００ＭＰａ以上を確保するため０．０５％以上とし、一方、０．８％を超えて添加するとベイナイト等の低温変態相を形成し、微細析出物による析出強化が不足し、強度が低下するため０．０５〜０．８％とする。より好ましくは０．１５〜０．４５％である。
【００４０】
Ｍｏは拡散速度が遅く、Ｔｉとともに析出する場合、析出物の成長速度が低下し、微細な析出物が得やすい。
【００４１】
（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）｝
本パラメータは、析出物の大きさに影響を与えるもので、０．５以上、１．５以下とした場合、粒径１０ｎｍ未満の微細析出物の形成が容易となり好ましい。微細なＴｉ，Ｍｏ系炭化物では、炭化物中のＴｉ、Ｍｏは原子比でＴｉ／Ｍｏが０．２〜２．０、更に微細な炭化物では０．７〜１．５であることが観察された。
【００４２】
更に、特性を向上させる場合、Ｎｂ、Ｖ、Ｗの一種または二種以上を添加することが好ましい。
【００４３】
Ｎｂ
ＮｂはＴｉと微細析出物を形成して強度向上に寄与する。また、組織を微細化し、結晶粒の整粒により延性を向上させる。０．０８％を超えると過度に微細化し、延性が低下するため０．０８％以下とする。より好ましくは０．０４％以下である。
【００４４】
Ｖ
ＶはＴｉと微細析出物を形成するが、０．１５％を超えると析出物が粗大化するようになるため、０．１５％以下とする。より好ましくは０．１０％以下である。
【００４５】
Ｗ
ＷはＴｉと微細析出物を形成するが、１．５％を超えると析出物が粗大化するようになるため、１．５％以下とする。より好ましくは１．０％以下である。
【００４６】
これらの元素の添加においては、Ｃ、Ｔｉ、Ｍｏ、Ｎｂ、Ｖ、Ｗの原子比を規定することが炭化物の微細化に有効である。
【００４７】
（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）＋（Ｖ／５１）＋（Ｗ／１９２）｝
本限定式は析出物の大きさに影響を与えるもので、（Ｃ／１２）／｛（Ｔｉ／４８）＋（Ｍｏ／９６）＋（Ｎｂ／９３）＋（Ｖ／５１）＋（Ｗ／１９２）｝を０．５以上、１．５以下とした場合、粒径１０ｎｍ未満の微細析出物の形成が容易となる。
【００４８】
Ｎｂ，Ｖ，Ｗの一種または二種以上を含む微細な炭化物の場合は、（Ｔｉ＋Ｎｂ＋Ｖ）／（Ｍｏ＋Ｗ）が０．２〜２．０、更に微細な炭化物の場合は０．７〜１．５であることが観察された。
【００４９】
また、本発明鋼では、部品成形時の被削性を向上させる場合は、Ｓ：０．０３〜０．１％とし、Ｐｂ≦０．２％、Ｃａ≦０．００５％，Ｂ≦０．０２％の一種または二種以上を添加することができる。
【００５０】
強度、延性を向上させる場合、Ｎｉ，Ｃｒの一種または二種をＮｉ≦２％、Ｃｒ≦２％の範囲で添加してもかまわない。
【００５１】
さらに、棒鋼の靭性を向上させる場合、不可避不純物であるＰ，Ｎを、Ｐ≦０．０４０％、Ｎ≦８０ｐｐｍに規制することが望ましい。
【００５２】
尚、これらの元素の含有量や添加の有無により本発明の効果が損なわれることはない。
【００５３】
３．製造条件
図１は本発明の熱間圧延棒鋼による非調質部品の概略製造工程図でＳ１は棒鋼製造工程、Ｓ２は搬送工程、Ｓ３は製品仕上げ過程を示す。棒線材製造工程（Ｓ１）で鋼塊を熱間圧延し棒鋼とし、製品仕上げ過程（Ｓ３）で棒鋼を熱間圧延加工し、所望の部品形状とした後、析出処理で微細析出物を析出させ引張強さ７００ＭＰａ以上とする。尚、本発明において、棒鋼の寸法によっては熱間圧延後の冷却速度を調整し、析出処理を省略することも可能である。
【００５４】
以下に望ましい製造工程について詳細に説明する。
【００５５】
圧延加熱温度
圧延加熱温度は１１００℃以上とする。本発明では、製品加工後の析出処理や圧延冷却過程で微細析出物を析出させるため、圧延時に溶解時から残存する炭化物を固溶させる。圧延加熱温度を１１００℃未満とした場合、溶解時から残存するＴｉ−Ｍｏ系炭化物等が固溶しないため１１００℃以上とする。
【００５６】
圧延後の冷却速度
冷却速度は、熱間圧延ままで微細析出物が析出するよう微細析出物の析出温度範囲の７００〜５５０℃を、微細析出物が得られる限界冷却速度（０．５℃／ｓｅｃ）以下の冷却速度で冷却する。この場合であっても、低Ｃ系組成により母相はフェライト単相組織となる。
【００５７】
析出処理
棒鋼の場合、その寸法によって、上述の冷却速度が得られない場合がある。そのような場合、製品成形後、析出処理を行う。析出処理では、母相をフェライト単相とし、強度向上に寄与する微細析出物を析出させることが必要で、加熱温度はベイナイトが生成しないよう５５０℃以上とし、７００℃を超えると析出物が粗大化するため５５０〜７００℃とする。
【００５８】
また、微細なＴｉ、Ｍｏなどの炭化物を生成、析出させるため該温度域において１０分以上保持する。
【００５９】
【実施例】
表１に示す組成の鋼を１５０ｋｇ真空溶解炉にて溶製し、１６０ｍｍ角に造塊後、ダミービレットに溶接し、種々の加熱温度、仕上げ温度でφ３０ｍｍ，φ１００ｍｍの棒鋼に圧延した。棒鋼の径は５５０〜７００℃の冷却速度を変化させるためφ３０ｍｍ、φ１００ｍｍとした。φ３０ｍｍの棒鋼については一部のものを除いて時効処理を行った。
【００６０】
供試鋼Ｎｏ．１〜１０は本発明例、Ｎｏ．１１〜１５は比較例で成分組成および／または製造条件が本発明の範囲外であり、Ｎｏ．１６は従来例でＳ４５Ｃ調質材の非調質鋼である。
【００６１】
それぞれの棒鋼について、組織観察を行い、降伏強さ、引張強さ、衝撃値を求めた。
【００６２】
引張試験はＪＩＳ４号試験片により常温での降伏強さ、引張強さを、衝撃試験はＪＩＳ３号のＵノッチ衝撃試験片により試験温度２０℃での吸収エネルギーを求めた。引張強さ７００ＭＰａ以上、降伏比０．８５以上が本発明例である。
【００６３】
組織観察は時効処理後、棒鋼断面を光学顕微鏡で観察し、また、透過型電子顕微鏡（ＴＥＭ）で薄膜観察を行った。析出物はエネルギー分散型Ｘ線分光装置（ＥＤＸ）により同定した。
【００６４】
表２に試験結果を示す。Ｎｏ．１〜１０はミクロ組織が本発明の規定を満足し、高強度、高靭性が得にくいとされるφ１００ｍｍのＮｏ，８〜１０を含めて７００ＭＰａ以上の引張強さ、０．８５以上の降伏比、１００Ｊ／ｃｍ^２以上の衝撃値であり、同強度の従来材（Ｎｏ．１６）に比較して高靭性、高降伏比であった。
【００６５】
一方、Ｎｏ．１１〜１５の比較例は、引張強さ、降伏比のいずれかまたは両者が本発明範囲外であった。
【００６６】
Ｎｏ．１１は圧延後の冷却速度が本発明範囲外で、時効処理を行わなかったため、微細析出物も観察されず、引張強さ、降伏比に劣る。
【００６７】
Ｎｏ．１２は圧延前の加熱温度が低く本発明範囲外で、引張強さ、降伏比に劣り、靭性も低い。低い加熱温度によりＴｉ，Ｍｏが十分固溶せず、時効処理においてこれら残存物を核に析出が生じたため、粗大析出物が形成されたためと思われる。
【００６８】
Ｎｏ．１３は時効処理温度が本発明範囲外で高く、析出物が粗大化し、降伏比が低い。
【００６９】
Ｎｏ．１４はＣ量が多く、限定式を満足せず析出物が粗大化し、降伏比が低い。
【００７０】
Ｎｏ．１５はＭｏ量、Ｔｉ量が本発明範囲外で、限定式を満足せず析出物が粗大化し、降伏比が低く、靭性も劣る。
【００７１】
【表１】

【００７２】
【表２】

【００７３】
【発明の効果】
本発明によれば、熱間圧延後、調質処理を行うことなく調質処理材と同等の強度、被削性を有する引張強さ７００ＭＰａ以上の熱間圧延非調質棒鋼およびその製造方法が得られ、産業上極めて有用である。
【図面の簡単な説明】
【図１】本発明鋼の製造工程の一例を示す図。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a hot-rolled non-heat treated steel bar and a method for producing the same, and in particular, even in a non-heat-treated heat-rolled steel sheet, has a tensile strength of 700 MPa or more, a yield ratio of 0.85 or more, high toughness, fatigue properties and It relates to one with excellent machinability.
[0002]
[Prior art]
Structural parts used in automobiles and construction machinery include tempered parts made by quenching and tempering carbon steel for machine structural use and alloy steel for machine structural use, and non-adjustable parts that secure strength by adjusting the composition and structure of the components without quenching and tempering. Quality parts are used.
[0003]
The steel used for non-heat treated parts generally has a ferrite / pearlite dual phase structure to which V and Nb are added. If the tensile strength is made comparable to that of the heat treated steel, the yield strength, the drawing value, and the impact value are reduced. On the other hand, it has been pointed out that when the yield strength is almost the same, the tensile strength, that is, the hardness is excessively increased and the machinability is reduced.
[0004]
JP-A-2001-123224 and JP-A-2001-131680 relate to a high-strength, high-yield-ratio, and high-toughness non-heat-treated steel, and have a structure having ferrite, bainitic ferrite, and pseudo-martensite. After the cold working, the steel is subjected to an aging treatment at a temperature of 600 ° C. or less to precipitate Cu and Ti—Nb-based carbides.
[0005]
[Problems to be solved by the invention]
However, it is not realistic to strictly control the ratio of a plurality of structures in actual production, and when utilizing precipitation strengthening by adding a large amount of Cu, it is necessary to add expensive Ni to prevent hot cracking. Not suitable as a structural part.
[0006]
Therefore, the present invention provides a hot-rolling type non-adjustable alloy which can obtain a strength of 700 Mpa or more, a yield ratio of 0.85 or more, a yield ratio and toughness with an inexpensive component composition without lowering the productivity even in actual machine production. It is an object of the present invention to provide a quality steel bar and a method for producing the same.
[0007]
[Means for Solving the Problems]
The present inventors have conducted various studies on steel bars which are non-heat treated and have strength and toughness comparable to those of the heat treated materials and do not impair the machinability, from the viewpoint of their component compositions and production conditions. It has been found that, when precipitation strengthening is performed with a large precipitate, the yield strength increases without excessively hardening the steel, and the toughness and machinability are not impaired.
[0008]
In the present invention, the hot-rolled non-heat-treated steel bar means a steel bar in which fine precipitates are precipitated by adjusting the cooling rate after the steel bar rolling, and the product does not require a tempering treatment after processing.
[0009]
The present invention has been made by further study based on the above findings, that is, the present invention,
1. Hot rolling having a ferrite single-phase structure, wherein fine precipitates having a grain size of less than 10 nm are dispersed and precipitated in the ferrite phase, and having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more. Mold non-heat treated steel bar.
[0010]
2. Steel composition in mass%, C ≦ 0.15%, Si ≦ 0.5%, Mn ≦ 2%, Al ≦ 0.1%, Ti: 0.03 to 0.35%, Mo: 0.05 2. A hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to 1 above, comprising -0.8%, the balance being Fe and unavoidable impurities.
[0011]
3. 3. The hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to 2, wherein the steel composition further satisfies the formula (1).

However, each element is the content (% by mass).
[0012]
4. 4. The hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to any one of 1 to 3, wherein the fine precipitates are carbides of Ti and Mo. .
[0013]
5. (2) The steel composition further contains one or more of Nb ≦ 0.08%, V ≦ 0.15%, and W ≦ 1.5% in mass%, and the tensile strength is 700 MPa or more and 0.85 or more. A hot-rolled non-heat treated steel bar having the above yield ratio.
[0014]
6. 6. A hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to 5, wherein the steel composition further satisfies the formula (2).
0.5 ≦ (C / 12) / ｛(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)｝ ≦ 1.5 (2)
However, each element is set to the content (%), and those not included are set to 0.
[0015]
7. 7. The heat having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to 5 or 6, wherein the fine precipitates are carbides containing Ti, Mo, Nb, V, and W. Cold rolled non-heat treated steel bar.
[0016]
8. The steel composition further includes one or more of S: 0.03 to 0.1%, Pb ≤ 0.2%, Ca ≤ 0.005%, and B ≤ 0.02% by mass%. The hot-rolled non-heat treated steel bar according to any one of 2, 3, 5, and 6, having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more.
[0017]
9. After heating the steel according to any one of 2, 3, 5, 6, and 8 to 1100 ° C. or more, hot-rolling the steel at a finish rolling temperature of 800 ° C. or more, and then cooling to 700 to 550 ° C. A hot rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more, characterized by cooling at a cooling rate of 0.5 ° C./sec or less.
[0018]
10. After heating the steel according to any one of 2, 3, 5, 6, and 8 to 1100 ° C. or more, hot-rolling the steel at a finish rolling temperature of 800 ° C. or more, and then cooling to 700 to 550 ° C. Hot rolling having a tensile strength of not less than 700 MPa and a yield ratio of not less than 0.85 characterized by cooling at a cooling rate of more than 0.5 ° C./sec and then reheating at 550 to 700 ° C. for not less than 10 minutes. Manufacturing method of mold non-heat treated steel bars.
[0019]
BEST MODE FOR CARRYING OUT THE INVENTION
The microstructure, component composition and production conditions of the present invention will be described in detail below.
[0020]
1. Microstructure The steel of the present invention defines its microstructure as a structure containing a fine precipitate having a grain size of less than 10 nm in a ferrite single phase structure.
[0021]
If a ferrite single phase structure, toughness comparable to the tempered material is obtained, and when fine precipitates are dispersed and precipitated in the structure, high yield strength comparable to the tempered material without impairing machinability can get. Fine precipitates are precipitated by adjusting the cooling rate after hot rolling or by performing a precipitation treatment after hot rolling.
[0022]
In the present invention, the fine precipitate has a particle size of less than 10 nm. If the particle size of the precipitate is 10 nm or more, the required tensile strength of 700 MPa or more as a component for mechanical structures of transportation equipment and construction equipment cannot be obtained.
[0023]
Further, when fine precipitates having a grain size of less than 10 nm are precipitated in the ferrite single phase structure, the yield ratio becomes 0.85 or more, and the increase in tensile strength is suppressed with respect to the increase in yield strength. Machinability comparable to tempered steel is obtained. If either the parent phase or the precipitate falls outside the scope of the present invention, the yield ratio will be less than 0.85.
[0024]
The smaller the particle size of the fine precipitates is, the more effective it is to improve the strength, preferably 5 nm, more preferably 3 nm or less. Such fine precipitates are carbides containing a complex of Ti and Mo, and also Nb, V, Carbides containing one or more of W are preferred.
[0025]
Although the distribution form of these fine precipitates is not particularly defined, it is desirable that they are uniformly dispersed (dispersed and precipitated) in the parent phase.
[0026]
In the present invention, if the size of the fine precipitates is 90% or more of the total precipitates, the desired tensile strength of 700 MPa or more can be obtained. However, a precipitate having a size of 10 nm or more consumes a precipitate-forming element and adversely affects the strength.
[0027]
Although the effect of the present invention is not impaired even if a small amount of Fe carbide is contained separately from the above-mentioned precipitates, the toughness is impaired when a large amount of Fe carbide having an average particle size of 1 μm or more impairs toughness. It is desirable that the upper limit of the size of the Fe carbide contained is 1 μm, and the content is 1% or less of the whole.
[0028]
The ratio of the fine precipitates to the total precipitates is determined by the following method. An electron microscope sample is prepared by an electrolytic polishing method using a twin jet method, and observed at an acceleration voltage of 200 kV. At this time, the crystal orientation of the parent phase is controlled so that the fine precipitate has a measurable contrast with respect to the parent phase, and the defocus is shifted from the normal focus to minimize the number of precipitates. Observe by method.
[0029]
In addition, the thickness of the sample in the region where the precipitation particles are measured is evaluated by measuring the elastic scattering peak and the inelastic scattering peak intensity using electron energy loss spectroscopy.
[0030]
According to this method, the measurement of the number of particles and the measurement of the sample thickness can be performed for the same region. The measurement of the number of particles and the particle diameter is performed on four portions of a 0.5 × 0.5 μm region of the sample, and the number of precipitates distributed per 1 μm ² is calculated as the number for each particle size.
[0031]
From this value and the sample thickness, the number of precipitates per particle size distributed per 1 μm ³ is calculated, and the ratio of the precipitates having a diameter of less than 10 nm to the measured total precipitates is calculated.
[0032]
In the present invention, the ferrite single-phase structure is defined as having a ferrite area ratio of 95% or more, and preferably 98% or more, when observed in a cross-sectional structure (observation with a 200-fold optical microscope).
[0033]
2. Ingredient Composition The steel of the present invention can achieve desired performance with the above-mentioned microstructure, but the following ingredient composition is preferred.
[0034]
C
If C is contained in excess of 0.15%, the fine precipitates become coarse and the strength is reduced, so that the content is preferably 0.15% or less. More preferably, it is 0.03% or more and 0.12% or less.
[0035]
Si
Si is added to improve cold workability. If it exceeds 0.5%, its effect will be impaired, so it is set to 0.5% or less. It is more preferably at most 0.15%.
[0036]
Mn
Mn is added to improve strength and ductility. If it exceeds 2%, its effect is impaired, so it is set to 2% or less. More preferably, it is 0.5% or more and 1.8% or less.
[0037]
Al
Al acts as a deoxidizing agent. In addition, N and AlN are formed to improve the quenching effect of B. If it exceeds 0.1%, the effect is saturated, so the content is made 0.1% or less. More preferably, it is 0.05% or less.
[0038]
Ti
Ti is added in order to finely precipitate a precipitate containing Ti-Mo-based carbide together with Ti-based carbide and Mo, and to improve the strength. To secure a tensile strength of 700 MPa or more, the content is made 0.03% or more. On the other hand, if added in excess of 0.35%, the precipitate becomes coarse and the strength and toughness are reduced, so that the content is made 0.03 to 0.35%. More preferably, it is 0.03 to 0.20%.
[0039]
Mo
Mo is added in order to finely precipitate a precipitate containing Ti-Mo-based carbide together with Mo-based carbide and Ti, and to improve the strength. In order to secure a tensile strength of 700 MPa or more, the content is made 0.05% or more. On the other hand, if added in excess of 0.8%, a low-temperature transformation phase such as bainite is formed, precipitation strengthening by fine precipitates is insufficient, and strength is reduced. Therefore, it is set to 0.05 to 0.8%. More preferably, it is 0.15 to 0.45%.
[0040]
Mo has a low diffusion rate, and when precipitated with Ti, the growth rate of the precipitates is reduced, and fine precipitates are easily obtained.
[0041]
(C / 12) / {(Ti / 48) + (Mo / 96)}
This parameter affects the size of the precipitate, and is preferably 0.5 or more and 1.5 or less, because it is easy to form a fine precipitate having a particle size of less than 10 nm, which is preferable. In fine Ti, Mo-based carbides, it was observed that Ti / Mo in the carbide had an atomic ratio of Ti / Mo of 0.2 to 2.0, and that of finer carbides was 0.7 to 1.5. .
[0042]
In order to further improve the characteristics, it is preferable to add one or more of Nb, V and W.
[0043]
Nb
Nb forms fine precipitates with Ti and contributes to improvement in strength. Further, the structure is refined, and the ductility is improved by sizing the crystal grains. If it exceeds 0.08%, it is excessively fine and the ductility decreases, so the content is made 0.08% or less. More preferably, it is 0.04% or less.
[0044]
V
V forms fine precipitates with Ti, but if it exceeds 0.15%, the precipitates become coarse, so that V is set to 0.15% or less. More preferably, it is 0.10% or less.
[0045]
W
W forms fine precipitates with Ti, but if it exceeds 1.5%, the precipitates become coarse, so the content is made 1.5% or less. More preferably, it is 1.0% or less.
[0046]
In addition of these elements, it is effective to define the atomic ratio of C, Ti, Mo, Nb, V, and W to make the carbide finer.
[0047]
(C / 12) / {(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 192)}
This limiting formula affects the size of the precipitate, and is (C / 12) / ｛(Ti / 48) + (Mo / 96) + (Nb / 93) + (V / 51) + (W / 192) When｝ is 0.5 or more and 1.5 or less, it becomes easy to form fine precipitates having a particle size of less than 10 nm.
[0048]
(Ti + Nb + V) / (Mo + W) is 0.2 to 2.0 for a fine carbide containing one or more of Nb, V, and W, and 0.7 to 1.5 for a finer carbide. Was observed.
[0049]
Further, in the steel of the present invention, when improving the machinability at the time of forming parts, S: 0.03 to 0.1%, Pb ≦ 0.2%, Ca ≦ 0.005%, B ≦ 0. One or more of 02% can be added.
[0050]
When improving the strength and ductility, one or two of Ni and Cr may be added in the range of Ni ≦ 2% and Cr ≦ 2%.
[0051]
Furthermore, when improving the toughness of a steel bar, it is desirable that P and N, which are inevitable impurities, be regulated to P ≦ 0.040% and N ≦ 80 ppm.
[0052]
The effects of the present invention are not impaired by the content of these elements or the presence or absence of these elements.
[0053]
3. 1. Manufacturing Conditions FIG. 1 is a schematic manufacturing process diagram of a non-heat treated part using the hot-rolled steel bar of the present invention. After the steel ingot is hot-rolled into a bar in the bar and wire manufacturing process (S1), the bar is hot-rolled in the product finishing process (S3) to obtain a desired part shape, and fine precipitates are precipitated by a precipitation process. The tensile strength is 700 MPa or more. In the present invention, depending on the size of the steel bar, the cooling rate after hot rolling can be adjusted, and the precipitation treatment can be omitted.
[0054]
Hereinafter, a desirable manufacturing process will be described in detail.
[0055]
Rolling heating temperature The rolling heating temperature is 1100 ° C. or higher. In the present invention, in order to precipitate fine precipitates in the precipitation treatment after the product processing or in the rolling and cooling process, the carbide remaining from the melting at the time of rolling is dissolved. When the rolling heating temperature is lower than 1100 ° C., the temperature is set to 1100 ° C. or higher because Ti-Mo-based carbide and the like remaining after melting do not form a solid solution.
[0056]
Cooling rate after rolling The cooling rate is set at 700 to 550 ° C., which is the precipitation temperature range of the fine precipitate, so that the fine precipitate is deposited as hot rolled. sec) Cool at the following cooling rate. Even in this case, the matrix has a ferrite single phase structure due to the low C-based composition.
[0057]
In the case of the precipitation-treated steel bar, the above-mentioned cooling rate may not be obtained depending on its size. In such a case, a precipitation treatment is performed after the product is formed. In the precipitation treatment, it is necessary to use a ferrite single phase as a matrix and precipitate fine precipitates contributing to the improvement of strength. 550-700 ° C.
[0058]
In order to generate and precipitate fine carbides such as Ti and Mo, the temperature is maintained for 10 minutes or more in the temperature range.
[0059]
【Example】
A steel having the composition shown in Table 1 was melted in a 150 kg vacuum melting furnace, ingoted into a square of 160 mm, then welded to a dummy billet, and rolled at various heating and finishing temperatures into steel bars of φ30 mm and φ100 mm. The diameter of the steel bar was φ30 mm and φ100 mm in order to change the cooling rate at 550 to 700 ° C. The aging treatment was performed on a steel bar of φ30 mm except for a part.
[0060]
Test steel No. Nos. 1 to 10 are examples of the present invention; Nos. 11 to 15 are comparative examples in which the component composition and / or the production conditions are out of the scope of the present invention. Reference numeral 16 denotes a conventional example, a non-heat-treated steel of S45C heat-treated material.
[0061]
The structure of each steel bar was observed, and the yield strength, tensile strength, and impact value were determined.
[0062]
In the tensile test, the yield strength and the tensile strength at room temperature were determined using a JIS No. 4 test piece, and in the impact test, the absorbed energy at a test temperature of 20 ° C. was obtained using a JIS No. 3 U-notch impact test piece. The present invention has a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more.
[0063]
After the aging treatment, the microstructure was observed by observing the cross section of the steel bar with an optical microscope, and the thin film was observed with a transmission electron microscope (TEM). The precipitate was identified by an energy dispersive X-ray spectrometer (EDX).
[0064]
Table 2 shows the test results. No. Nos. 1 to 10 have a tensile strength of 700 MPa or more, including φ100 mm Nos. 8 to 10 whose microstructure satisfies the requirements of the present invention and high strength and high toughness are difficult to obtain, and a yield ratio of 0.85 or more. , 100 J / cm ² or more, and high toughness and high yield ratio as compared with the conventional material (No. 16) having the same strength.
[0065]
On the other hand, No. In Comparative Examples 11 to 15, one or both of the tensile strength and the yield ratio were out of the range of the present invention.
[0066]
No. In No. 11, since the cooling rate after rolling was out of the range of the present invention and the aging treatment was not performed, fine precipitates were not observed, and the tensile strength and the yield ratio were poor.
[0067]
No. No. 12 has a low heating temperature before rolling and is out of the range of the present invention, and is inferior in tensile strength and yield ratio and low in toughness. It is considered that Ti and Mo did not sufficiently form a solid solution due to the low heating temperature, and these residues were deposited as nuclei during the aging treatment, so that coarse precipitates were formed.
[0068]
No. In No. 13, the aging temperature is high outside the range of the present invention, the precipitate is coarsened, and the yield ratio is low.
[0069]
No. No. 14 has a large C content, does not satisfy the limiting formula, coarsens the precipitate, and has a low yield ratio.
[0070]
No. In No. 15, the Mo content and the Ti content are out of the range of the present invention, do not satisfy the limiting formula, the precipitates are coarsened, the yield ratio is low, and the toughness is poor.
[0071]
[Table 1]

[0072]
[Table 2]

[0073]
【The invention's effect】
According to the present invention, a hot-rolled non-heat-treated steel bar having a tensile strength of 700 MPa or more having a strength equivalent to that of a heat-treated material without heat treatment and a machinability after hot rolling, and a method for producing the same. Obtained and extremely useful industrially.
[Brief description of the drawings]
FIG. 1 is a diagram showing an example of a production process of the steel of the present invention.

Claims (10)

Hot rolling having a ferrite single-phase structure and fine precipitates having a grain size of less than 10 nm dispersed and precipitated in the ferrite phase and having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more. Mold non-heat treated steel bar. Steel composition in mass%, C ≦ 0.15%, Si ≦ 0.5%, Mn ≦ 2%, Al ≦ 0.1%, Ti: 0.03 to 0.35%, Mo: 0.05 The hot-rolled non-heat-treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to claim 1, comprising 0.8% to 0.8%, the balance being Fe and unavoidable impurities. The hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to claim 2, wherein the steel composition further satisfies the formula (1).

However, each element is the content (% by mass). 4. The hot-rolling mold as claimed in claim 1, wherein the fine precipitates are carbides of Ti and Mo and have a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more. Quality steel bars. 3. The steel composition according to claim 2, further comprising one or more of Nb ≦ 0.08%, V ≦ 0.15%, and W ≦ 1.5% in mass%. A hot-rolled non-heat treated steel bar having a yield ratio of .85 or more. The hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more according to claim 5, wherein the steel composition further satisfies the formula (2).
0.5 ≦ (C / 12) / ｛(Ti / 48) + (Mo / 96) + (Nb / 93)
+ (V / 51) + (W / 184)｝ ≦ 1.5 (2)
However, each element is set to the content (%), and those not included are set to 0. The fine precipitate is a carbide containing Ti, Mo, Nb, V and W, and has a tensile strength of not less than 0.85 at a tensile strength of 700 MPa or more according to claim 5 or 6. Hot rolled non-heat treated steel bars. The steel composition further includes one or more of S: 0.03 to 0.1%, Pb ≤ 0.2%, Ca ≤ 0.005%, and B ≤ 0.02% by mass%. The hot-rolled non-heat-treated steel bar according to any one of claims 2, 3, 5, and 6, having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more. After heating the steel according to any one of claims 2, 3, 5, 6, and 8 to 1100 ° C. or more, hot-rolling the steel at a finish rolling temperature of 800 ° C. or more, and then cooling to 700 to 700 ° C. A method for producing a hot-rolled non-heat treated steel bar having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more, characterized by cooling 550 ° C. at a cooling rate of 0.5 ° C./sec or less. After heating the steel according to any one of claims 2, 3, 5, 6, and 8 to 1100 ° C. or more, hot-rolling the steel at a finish rolling temperature of 800 ° C. or more, and then cooling to 700 to 700 ° C. Heat having a tensile strength of 700 MPa or more and a yield ratio of 0.85 or more characterized by cooling 550 ° C. at a cooling rate of more than 0.5 ° C./sec, and then reheating at 550 to 700 ° C. for 10 minutes or more. A method for producing a cold-rolled non-heat treated steel bar.

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