JP2004035948A - High-strength high-stiffness steel and manufacturing method therefor - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a high-strength high-stiffness steel which can be formed with hot working, has a Young's modulus of 240 GPa or higher, and is manufactured with a smelting method. <P>SOLUTION: The high-strength high-stiffness steel manufactured by casting of the molten steel has hexagonal borides and cubic carbonitrides with particle diameters of 20 μm or less, dispersed in a matrix of a Fe-based alloy at a volume fraction of 10-35% in total. It is further preferable that the hexagonal boride is one or more compounds of TiB<SB>2</SB>, VB<SB>2</SB>, NbB<SB>2</SB>and a composite boride thereof, and that the cubic carbonitride is one or more compounds of Ti (C, N), V (C, N), Nb (C, N), Ta (C, N), Hf (C, N) and a composite carbonitride thereof. <P>COPYRIGHT: (C)2004,JPO

Description

但し、各式中の元素記号は、その元素の質量％での含有量を表す。
【００１７】
（４）質量％で、Ｃ：０．２〜１．７％、Ｓｉ：０．０５〜０．５％、Ｍｎ：０．２〜１．５％、Ｔｉ：１〜１２％、Ｍｏ：０．１〜２％、Ｂ：０．５〜３．２％、Ｎ：０．００１〜０．０１％、Ｖ：０〜１２％、Ｎｂ：０〜１２％、Ｔａ：０〜１％、Ｈｆ：０〜１％、Ｃａ：０〜０．００７％、Ｍｇ：０〜０．００７％及びＮｄ：０〜０．００７％を含有し、残部がＦｅ及び不純物からなり、前記▲１▼式で表されるＳＢが１００を超えて１０５０未満、前記▲２▼式で表されるＳＣが５０を超えて５５０未満である鋼における各元素の化合物を鋼中に完全に溶解させた溶鋼の鋳造過程で、先ず六方晶系硼化物又は立方晶系炭窒化物を晶出させ、次いで、残留溶鋼から六方晶系硼化物、立方晶系炭窒化物及びＦｅマトリックス相の３相を晶出させて鋼塊とし、この鋼塊を３００℃以下の温度域まで冷却し、更に下記▲４▼式で表されるＨＴ（℃）以下の温度で均質化処理を行った後、ＨＴ（℃）以下の温度で総減面率が５０％以上となる熱間加工を行い、この後更に、８５０〜１０５０℃で焼ならしを行って３００℃以下まで冷却し、次いで、７００℃以下で焼戻し処理することを特徴とする高強度高剛性鋼の製造方法。
【００１８】

但し、▲４▼式中の元素記号は、その元素の質量％での含有量を表す。
【００１９】
本発明でいう「粒径」とは、具体的には個々の粒子の短径と長径の和の１／２で定義される値をいう。
【００２０】
本発明で規定する六方晶系硼化物及び立方晶系炭窒化物は、例えば、鋼中の晶出物又は析出物を電解抽出法等の通常の方法によって抽出し、得られた残渣をＸ線回折することによって確認することができる。
【００２１】
ＴｉＢ_２　、ＶＢ_２　及びＮｂＢ_２　の複合硼化物とは、「Ｘ」をＴｉ、Ｖ及びＮｂから選択されるいずれか２種以上として「ＸＢ_２　」で表される硼化物、具体的には、例えば、（Ｔｉ、Ｖ）Ｂ_２　、（Ｖ、Ｎｂ）Ｂ_２　や（Ｔｉ、Ｖ、Ｎｂ）Ｂ_２　などを指す。又、Ｔｉ（Ｃ、Ｎ）、Ｖ（Ｃ、Ｎ）、Ｎｂ（Ｃ、Ｎ）、Ｔａ（Ｃ、Ｎ）及びＨｆ（Ｃ、Ｎ）の複合炭窒化物とは、「Ｘ」をＴｉ、Ｖ、Ｎｂ、Ｔａ及びＨｆから選択されるいずれか２種以上として「Ｘ（Ｃ、Ｎ）」で表される炭窒化物、具体的には、例えば、（Ｔｉ、Ｖ）（Ｃ、Ｎ）、（Ｖ、Ｎｂ、Ｈｆ）（Ｃ、Ｎ）、（Ｔｉ、Ｖ、Ｎｂ、Ｈｆ）（Ｃ、Ｎ）や（Ｔｉ、Ｖ、Ｎｂ、Ｈｆ、Ｔａ）（Ｃ、Ｎ）などを指す。
【００２２】
なお、上記の六方晶系硼化物及び立方晶系炭窒化物は、その生成形態から下記（イ）〜（ハ）の３種類に分けられる。
【００２３】
（イ）初晶型の六方晶系硼化物及び立方晶系炭窒化物：
鋳造の冷却過程で溶鋼から初めに晶出するものであって、その粒径は１μｍ以上である。
【００２４】
（ロ）共晶型の六方晶系硼化物及び立方晶系炭窒化物：
上記（イ）の初晶型の硼化物又は炭窒化物が晶出した後、残留溶鋼から硼化物、炭化物及びＦｅマトリックスがほぼ同時に晶出して得られるものであって、その粒径は１μｍ未満である。
【００２５】
（ハ）析出型の六方晶系硼化物及び立方晶系炭窒化物：
鋼が完全に凝固した後の冷却過程、又は凝固後に施される熱処理等の過程で形成されるものであって、その粒径は０．１μｍ以下である。
【００２６】
ここで、上記（イ）と（ロ）の晶出型の六方晶系硼化物及び立方晶系炭窒化物は、光学顕微鏡や走査型電子顕微鏡を用いて観察できるし、（ハ）の析出型の六方晶系硼化物及び立方晶系炭窒化物は、例えば、加速電圧が１００〜２００ｋＶの透過電子顕微鏡を用いることにより観察することができる。
【００２７】
本発明者らは、溶製法によって製造した場合にも安定して微細且つ均一に晶出又は析出し、鋳造ままの材料だけではなく、鋳造後に熱間加工を加えても高いヤング率を確保させることができる分散粒子を見出すために、鋼の化学組成に加えて、鋳造条件、熱処理条件や熱間加工といった製造条件を種々変えて検討を行った。その結果、下記（ａ）〜（ｉ）の知見を得た。
【００２８】
（ａ）溶鋼の鋳造時に生成する化合物のうち、高いヤング率を有するのは、六方晶系硼化物及び立方晶系炭窒化物である。すなわち溶製鋼の場合には、上記の六方晶系硼化物及び立方晶系炭窒化物以外の化合物は、その化合物中にＦｅが分配され、化合物自体のヤング率が低下してしまう。これに対し、溶製鋼の場合であっても六方晶系硼化物及び立方晶系炭窒化物中にＦｅはほとんど分配されないので、高いヤング率が確保される。したがって、六方晶系硼化物及び立方晶系炭窒化物は溶製鋼の高ヤング率化に有効である。
【００２９】
（ｂ）六方晶系硼化物のうちでも、「Ｘ」をＴｉ、Ｖ及びＮｂから選択される１種以上として、「ＸＢ_２　」で表される六方晶系硼化物が最も高いヤング率を有し、溶製鋼を高ヤング率化するのに適している。
【００３０】
（ｃ）立方晶系炭窒化物のうちでも、「Ｘ」をＴｉ、Ｖ、Ｎｂ、Ｔａ及びＨｆから選択される１種以上として、「Ｘ（Ｃ、Ｎ）」で表される立方晶系炭窒化物が最も高いヤング率を有し、溶製鋼を高ヤング率化するのに有効である。
【００３１】
（ｄ）析出により得られた立方晶系炭窒化物は、強力な析出強化作用も有する。
【００３２】
（ｅ）六方晶系硼化物と立方晶系炭窒化物のいずれか一方で溶製鋼の高ヤング率化を実現するためには、前者の場合にはＢを後者の場合にはＣを、それぞれ多量の合金元素とともに添加する必要があるが、これらの場合には、溶鋼を鋳造する際の冷却過程で、溶鋼中に２０μｍを超える粒径の粗大な六方晶系硼化物又は立方晶系炭窒化物が晶出し、更にこれらが溶鋼中で凝集粗大化する。このような粒径が２０μｍを超える粗大な六方晶系硼化物又は立方晶系炭窒化物はＦｅマトリックスの界面との密着性が低いので、鋼のヤング率が低下してしまう。更に、鋳造後、熱間加工などの塑性変形を加えて成形すると、六方晶系硼化物又は立方晶系炭窒化物とＦｅマトリックス間に欠陥が生じて、ヤング率の一層の低下が生じる。このため、上記の場合には、鋳造ままの材料や熱間加工材において粒子の体積分率が１０％を超えると、ヤング率の複合則が成立せず、例えば２４０ＧＰａという高いヤング率を実現させることが困難である。
【００３３】
（ｆ）ＣとＢの双方を含む溶鋼から六方晶系硼化物と立方晶系炭窒化物のいずれか一方又は双方が晶出する場合、六方晶系硼化物に対してはＣが、又、立方晶系炭窒化物に対してはＢが、それぞれ界面エネルギーを低下させるので、初晶の六方晶系硼化物又は／及び初晶の立方晶系炭窒化物の凝集粗大化が抑制される。更に、初晶の六方晶系硼化物又は／及び立方晶系炭窒化物が晶出した後、低温の残留溶鋼から六方晶系硼化物、立方晶系炭窒化物及びＦｅマトリックス相をほぼ同時に晶出（共晶）させることにより、晶出粒子の粒径は２０μｍ以下となる。その結果、鋳造ままの材料だけではなく、その後に熱間加工を加えてもヤング率は低下せず、粒子の体積分率が１０〜３５％の範囲でもヤング率の複合則が成立する。
【００３４】
（ｇ）上記（ｆ）に記載の凝固形態を、鋳造時の通常の冷却条件、例えば、０．０５〜１℃／秒の冷却速度範囲で実現するためには、六方晶系硼化物と溶鋼が共存する温度範囲（固液共存域）及び、立方晶系炭窒化物と溶鋼が共存する温度範囲（固液共存域）を各々下記のようにすればよい。
【００３５】
・１００℃＜溶鋼と六方晶系硼化物の共存温度域＜１０５０℃、
・５０℃＜溶鋼と立方晶系炭窒化物の共存温度域＜５５０℃。
【００３６】
そこで、上記の温度範囲について、示差熱分析や、各温度域からの凍結組織を観察し、成分元素の影響を調査した。その結果、上記の溶鋼と六方晶系硼化物の共存温度域及び溶鋼と立方晶系炭窒化物の共存温度域はそれぞれ下記のＳＢ指数及びＳＣ指数で表記できることが明らかとなった。
【００３７】
すなわち、上記（ｆ）に記載の凝固形態を実現するためには、各式中の元素記号をその元素の質量％での含有量として、下記▲１▼式で表されるＳＢが１００を超えて１０５０未満、下記▲２▼式で表されるＳＣが５０を超えて５５０未満となればよい。
【００３８】

【００３９】
（ｈ）特定の化学組成を有し、且つ、前記（ｇ）の条件を満足する場合には、オーステナイト安定域で熱間加工を行えば、ヤング率の低下を最小限に抑えることができる。オーステナイト域の上限ＨＴ（℃）は式中の元素記号をその元素の質量％での含有量として前記▲４▼の実験式で表すことができ、したがって、ＨＴ（℃）以下の温度で熱間加工を行えばよい。
【００４０】
（ｉ）７００℃以下の温度で焼戻し処理することによって、微細炭化物が析出し、鋼を高強度化することができる。
【００４１】
本発明は、上記の知見に基づいて完成されたものである。
【００４２】
【発明の実施の形態】
以下、本発明の各要件について詳しく説明する。
【００４３】
（Ａ）Ｆｅ基合金のマトリックス中の粒子
ヤング率を高め、鋼に２４０ＧＰａ以上のヤング率を確保させるためには、六方晶系硼化物及び立方晶系炭窒化物をＦｅ基合金のマトリックス中に体積分率で合計１０％以上分散させなければならない。一方、Ｆｅ基合金のマトリックス中に分散する六方晶系硼化物及び立方晶系炭窒化物が体積分率で合計３５％を超えると、ヤング率が低下する上に、鋼塊に割れやポロシティーが発生し、その後の加工が困難となる。したがって、本発明においては、Ｆｅ基合金のマトリックス中に六方晶系硼化物及び立方晶系炭窒化物を体積分率で合計１０〜３５％分散させることとした。
【００４４】
本発明で規定する六方晶系硼化物及び立方晶系炭窒化物には、既に述べた粒径が１μｍ以上の初晶型の六方晶系硼化物及び立方晶系炭窒化物、粒径が１μｍ未満の共晶型の六方晶系硼化物及び立方晶系炭窒化物、及び、粒径が０．１μｍ以下の析出型の六方晶系硼化物及び立方晶系炭窒化物のいずれをも含む。
【００４５】
ここで、或る相の組成、結晶構造及び体積割合は、鋼中のＦｅマトリックス以外の粒子を電解抽出し、得られた残渣の定量化学分析、Ｘ線回折強度測定、個々の粒子のＥＤＸ分析及び鋼の密度測定によって求めることができる。
【００４６】
前記した晶出型及び析出型の六方晶系硼化物及び立方晶系炭窒化物は、光学顕微鏡、走査型電子顕微鏡及び、例えば、加速電圧が１００〜２００ｋＶの透過電子顕微鏡を用いて観察することができるので、観察によって得られた像を画像解析して短径と長径を測定し、その和の１／２から各粒子の粒径を求めることができる。六方晶系硼化物及び立方晶系炭窒化物に関し、上記のようにして求めた粒径が２０μｍを超える場合には、所望の２４０ＧＰａ以上の高いヤング率を安定して確保することが困難となる。したがって、Ｆｅ基合金のマトリックス中に体積分率で合計１０〜３５％分散する六方晶系硼化物及び立方晶系炭窒化物の粒径を２０μｍ以下とした。なお、上記六方晶系硼化物及び立方晶系炭窒化物の粒径の下限値は規定する必要はないが、加速電圧が１００〜２００ｋＶの透過電子顕微鏡で観察できる５ｎｍ程度を下限値としてもよい。
【００４７】
なお、Ｆｅ基合金のマトリックス中に体積分率で合計１０〜３５％分散する六方晶系硼化物と立方晶系炭窒化物の割合は、５：１〜１：１であることが好ましい。この条件を満たすことで極めて健全な鋼塊が得られる。
【００４８】
Ｆｅ基合金のマトリックス中に体積分率で立方晶系炭窒化物とともに合計で１０〜３５％分散させる粒子としての六方晶系硼化物は、ＴｉＢ_２　、ＶＢ_２　、ＮｂＢ_２　及びこれらの複合硼化物のうちの１種以上であることが好ましい。なお、ＴｉＢ_２　、ＶＢ_２　及びＮｂＢ_２　の複合硼化物とは、「Ｘ」をＴｉ、Ｖ及びＮｂから選択されるいずれか２種以上として「ＸＢ_２　」で表される硼化物、具体的には、例えば、（Ｔｉ、Ｖ）Ｂ_２　、（Ｖ、Ｎｂ）Ｂ_２　や（Ｔｉ、Ｖ、Ｎｂ）Ｂ_２　などを指すことは既に述べたとおりである。
【００４９】
更に、Ｆｅ基合金のマトリックス中に体積分率で六方晶系硼化物とともに合計１０〜３５％分散させる粒子としての立方晶系炭窒化物は、Ｔｉ（Ｃ、Ｎ）、Ｖ（Ｃ、Ｎ）、Ｎｂ（Ｃ、Ｎ）、Ｔａ（Ｃ、Ｎ）、Ｈｆ（Ｃ、Ｎ）及びこれらの複合炭窒化物のうちの１種以上であることが好ましい。なお、既に述べたように、Ｔｉ（Ｃ、Ｎ）、Ｖ（Ｃ、Ｎ）、Ｎｂ（Ｃ、Ｎ）、Ｔａ（Ｃ、Ｎ）又はＨｆ（Ｃ、Ｎ）の複合炭窒化物とは、「Ｘ」をＴｉ、Ｖ、Ｎｂ、Ｔａ及びＨａから選択されるいずれか２種以上として「Ｘ（Ｃ、Ｎ）」で表される炭窒化物、具体的には、例えば、（Ｔｉ、Ｖ）（Ｃ、Ｎ）、（Ｖ、Ｎｂ、Ｈｆ）（Ｃ、Ｎ）、（Ｔｉ、Ｖ、Ｎｂ、Ｈｆ）（Ｃ、Ｎ）や（Ｔｉ、Ｖ、Ｎｂ、Ｈｆ、Ｔａ）（Ｃ、Ｎ）などを指す。
【００５０】
（Ｂ）鋼の化学組成
次に、本発明に係る高強度高剛性鋼の好ましい化学組成に関して説明する。なお、各元素の含有量の「％」表示は「質量％」を意味する。
【００５１】
Ｃ：０．２〜１．７％
Ｃは、立方晶系炭窒化物を形成し、鋼のヤング率及び強度を高めるのに有効な元素である。更に、溶鋼中で晶出した硼化物の凝集粗大化を抑制する作用を有する。しかし、その含有量が０．２％未満の場合には、立方晶系炭窒化物の晶出量が少なく、所望の２４０ＧＰａ以上のヤング率が得られない場合がある。一方、その含有量が１．７％を超えると、晶出した炭窒化物が粗大化しやすくなって、所望のヤング率が得られない上に、鋳造性や熱間加工性に悪影響を及ぼす場合がある。したがって、Ｃの含有量は０．２〜１．７％とするのがよい。Ｃの含有量は、０．３〜１．２％であれば一層好ましく、０．４〜０．８％であれば極めて好ましい。
【００５２】
Ｓｉ：０．０５〜０．５％
Ｓｉは、脱酸に有効な元素であるが、その含有量が０．０５％未満では効果が得難い場合がある。一方、鋼がＴｉ、Ｖ及びＮｂなどを多量に含むと、粒界にラーベス相などの金属間化合物が生成しやすくなり、特に、Ｓｉの含有量が０．５％を超えると粒界金属間化合物が粗大化する場合がある。ラーベス相などの金属間化合物は熱間加工性の大きな低下をもたらす上に、ラーベス相自体が低ヤング率粒子であることから鋼のヤング率が低下してしまう。したがって、Ｓｉの含有量は０．０５〜０．５％とするのがよい。Ｓｉの含有量の上限は０．３％であることが更に望ましく、０．２％であれば一層よい。
【００５３】
Ｍｎ：０．２〜１．５％
本発明に係る高強度高剛性鋼は、特定のオーステナイト領域で均質化処理を行うことにより、加工性、強度が高まる。このため、オーステナイト安定化元素であるＭｎを用いてオーステナイト域を広げるのがよい。Ｍｎには、鋼の不純物であるＳを固定して、熱間加工性を高める作用もある。しかし、Ｍｎの含有量が０．２％未満では前記の効果が得られない場合がある。一方、その含有量が１．５％を超えると、ヤング率を低下させるＭ_２３Ｃ_６　型の炭化物が生成しやすくなる。したがって、Ｍｎ含有量は０．２〜１．５％とするのがよい。Ｍｎの含有量は、０．８〜１．２％であれば一層好ましい。
【００５４】
Ｔｉ：１〜１２％
Ｔｉは、ＴｉＢ_２　、（Ｔｉ、Ｖ）Ｂ_２　等の六方晶系硼化物、及びＴｉ（Ｃ、Ｎ）等の立方晶系炭窒化物を形成して鋼のヤング率を高めるのに有効な元素である。更に、鋼の高強度化にも有効である。しかし、その含有量が１％未満の場合には、前記の効果が得られない場合がある。一方、その含有量が１２％を超えると、六方晶系硼化物及び立方晶系炭窒化物が粗大化し、ヤング率が低下する場合がある。したがって、Ｔｉの含有量は１〜１２％とするのがよい。Ｔｉの含有量は、３〜１０％であれば一層好ましく、５〜９％であればより一層好ましく、７〜８％であれば極めて好ましい。
【００５５】
Ｍｏ：０．１〜２％
Ｍｏは、マトリックス中に固溶した状態で存在すると、粒界を強化して熱間加工性を著しく改善する作用を有する。更に、鋼の強度を高める作用も有する。しかし、その含有量が０．１％未満の場合には、前記の効果が得られない場合がある。一方、Ｍｏを２％を超えて含有させても、その効果は飽和する場合がありコストが嵩んでしまう。したがって、Ｍｏの含有量は０．１〜２％とするのがよい。Ｍｏの含有量のより望ましい範囲は０．２〜１％であり、０．３〜０．７％であれば極めて好ましい。
【００５６】
Ｂ：０．５〜３．２％
Ｂは、六方晶系硼化物を形成し、鋼のヤング率を高めるのに有効な元素である。しかし、その含有量が０．５％未満の場合には、形成される六方晶系硼化物の量が少なく、所望の２４０ＧＰａ以上のヤング率が得られない場合がある。一方、その含有量が３．２％を超えると、晶出する硼化物が粗大化し、ヤング率が低下する上に、熱間加工性に悪影響を及ぼす場合がある。したがって、Ｂの含有量は０．５〜３．２％とするのがよい。Ｂの含有量は１．０〜３．０％であれば一層好ましく、１．８〜２．７％であればより一層好ましく、２．０〜２．４％であれば極めて好ましい。
【００５７】
Ｎ：０．００１〜０．０１％
Ｎは、立方晶系炭窒化物を形成し、鋼のヤング率を高めるのに有効な元素である。しかし、その含有量が０．００１％未満の場合には、前記の効果が得られない場合がある。一方、その含有量が０．０１％を超えると、熱間加工性の低下を招く場合がある。したがって、Ｎの含有量は０．００１〜０．０１％とするのがよい。Ｎ含有量のより望ましい範囲は、０．００３〜０．００７％である。
【００５８】
本発明に係る高強度高剛性鋼は、上記のＣからＮまでの元素、Ｆｅ及び不純物以外に、更に、Ｎｂ、Ｖ、Ｈｆ及びＴａから選択される１種以上、又は／及びＣａ、Ｍｇ及びＮｄから選択される１種以上を選択的に含有していてもよい。
【００５９】
Ｖ：０〜１２％
Ｖは、添加すれば、ＶＢ_２　、（Ｔｉ、Ｖ）Ｂ_２　等の六方晶系硼化物、及びＶ（Ｃ、Ｎ）、（Ｔｉ、Ｖ）（Ｃ、Ｎ）等の立方晶系炭窒化物を形成して鋼のヤング率を高める作用を有する。又、鋼の強度を高める作用も有する。これらの効果を確実に得るには、Ｖは０．５％以上の含有量とすることが好ましい。しかし、その含有量が１２％を超えると加工性の著しい低下を招く場合がある。したがって、Ｖの含有量は０〜１２％とするのがよい。
【００６０】
Ｎｂ：０〜１２％
Ｎｂは、添加すれば、ＮｂＢ_２　、（Ｔｉ、Ｎｂ）Ｂ_２　等の六方晶系硼化物、及びＮｂ（Ｃ、Ｎ）、（Ｔｉ、Ｎｂ）（Ｃ、Ｎ）等の立方晶系炭窒化物を形成して鋼のヤング率を高める作用を有する。又、熱間加工性を高める作用や鋼の強度を高める作用も有する。これらの効果を確実に得るには、Ｎｂは０．０５％以上の含有量とすることが好ましい。しかし、その含有量が１２％を超えると加工性の著しい低下を招く場合がある。したがって、Ｎｂの含有量は０〜１２％とするのがよい。
【００６１】
Ｔａ：０〜１％
Ｔａは、添加すれば、Ｔａ（Ｃ、Ｎ）、（Ｔｉ、Ｔａ）（Ｃ、Ｎ）等の立方晶系炭窒化物を形成して鋼のヤング率を高める作用を有する。上記立方晶系炭窒化物は、凝固組織の微細化及び結晶粒の微細化作用を有し、こうした作用を通じて粒界偏析を低減させるので、熱間加工性を高める作用もある。更に、鋼の高強度化にも有効である。これらの効果を確実に得るには、Ｔａは０．０５％以上の含有量とすることが好ましい。しかし、その含有量が１％を超えると加工性の低下を招く場合がある。したがって、Ｔａの含有量は０〜１％とするのがよい。
【００６２】
Ｈｆ：０〜１％
Ｈｆは、添加すれば、Ｈｆ（Ｃ、Ｎ）、（Ｔｉ、Ｈｆ）（Ｃ、Ｎ）等の立方晶系炭窒化物を形成して鋼のヤング率を高める作用を有する。上記立方晶系炭窒化物は、凝固組織の微細化及び結晶粒の微細化作用を有し、こうした作用を通じて粒界偏析を低減させるので、熱間加工性を高める作用もある。更に、鋼の高強度化にも有効である。これらの効果を確実に得るには、Ｈｆは０．２％以上の含有量とすることが好ましい。しかし、その含有量が１％を超えると加工性の低下を招く場合がある。したがって、Ｈｆの含有量は０〜１％とするのがよい。
Ｃａ：０〜０．００７％
Ｃａは、添加すれば、鋼の不純物であるＳ及びＯ（酸素）を固定して、鋳造性や熱間加工性を高める作用を有する。この効果を確実に得るには、Ｃａは０．００２％以上の含有量とすることが好ましい。しかし、その含有量が０．００７％を超えると介在物が粗大化して熱間加工性の低下を招く場合がある。したがって、Ｃａの含有量は０〜０．００７％とするのがよい。
Ｍｇ：０〜０．００７％
Ｍｇは、添加すれば、鋼の不純物であるＳ及びＯ（酸素）を固定して、鋳造性や熱間加工性を高める作用を有する。この効果を確実に得るには、Ｍｇは０．００２％以上の含有量とすることが好ましい。しかし、その含有量が０．００７％を超えると介在物が粗大化して熱間加工性の低下を招く場合がある。したがって、Ｍｇの含有量は０〜０．００７％とするのがよい。
Ｎｄ：０〜０．００７％
Ｎｄは、添加すれば、鋼の不純物であるＳ及びＯ（酸素）を固定して、鋳造性や熱間加工性を高める作用を有する。この効果を確実に得るには、Ｎｄは０．００２％以上の含有量とすることが好ましい。しかし、その含有量が０．００７％を超えると介在物が粗大化して熱間加工性の低下を招く場合がある。したがって、Ｎｄの含有量は０〜０．００７％とするのがよい。
（Ｃ）鋼の製造方法
次に、本発明に係る高強度高剛性鋼の製造条件の一例について説明する。
【００６３】
（Ｃ−１）溶製条件
本発明に係る高強度高剛性鋼の組織を得るためには、前記（Ｂ）項で述べた化学組成に加えて、溶鋼と六方晶系硼化物の共存温度域に係る前記▲１▼式で表されるＳＢ及び、溶鋼と立方晶系炭窒化物の共存温度域に係る前記▲２▼式で表されるＳＣが、それぞれ、
１００＜ＳＢ＜１０５０、
５０＜ＳＣ＜５５０、
を満たす合金組成を選択し、各元素の化合物を鋼中に完全に溶解させた後、適正な凝固形態が得られるように鋳造することが望ましい。
【００６４】
なお、ＳＢの値が１０５０以上になると、初晶の六方晶系硼化物の粒径が２０μｍを超えて、鋳造ままの材料又は熱間加工材のヤング率の低下をもたらす場合がある。一方、ＳＢの値が１００以下になると、溶鋼から晶出する六方晶系硼化物の量が不十分となって、所望の２４０ＧＰａ以上のヤング率が得られない場合がある。
【００６５】
又、ＳＣの値が５５０以上になると、初晶の立方晶系炭窒化物の粒径が２０μｍを超えて、鋳造ままの材料又は熱間加工材のヤング率の低下をもたらす場合がある。一方、ＳＣの値が５０以下になると、溶鋼から晶出する立方晶系炭窒化物の量が不十分となって、所望の２４０ＧＰａ以上のヤング率が得られない場合がある。
【００６６】
（Ｃ−２）鋳造条件
前記（Ｂ）項で述べた化学組成に加えて、▲１▼式及び▲２▼式で表されるＳＢとＳＣが
１００＜ＳＢ＜１０５０、
及び
５０＜ＳＣ＜５５０、
を満たす鋼における各元素の化合物を鋼中に完全に溶解させた後の鋳造時の冷却過程で、先ず、六方晶系硼化物と立方晶系炭窒化物のいずれか一方又は双方を晶出させ、次いで、残留溶鋼から六方晶系硼化物、立方晶系炭窒化物及びＦｅマトリックス相の３相をほぼ同時に晶出（共晶）させることが望ましい。初晶をδフェライトなどのＦｅマトリックスとする場合には、ヤング率の低い斜方晶の硼化物やＦｅ含有量の多い炭化物が安定晶出又は析出して、本発明で規定する六方晶系硼化物や立方晶系炭窒化物が十分晶出しないことになって、所望の２４０ＧＰａ以上のヤング率が得られない場合がある。又、六方晶系硼化物、立方晶系炭窒化物及びＦｅマトリックス相の３相のほぼ同時の晶出（共晶）過程を経なければ、粒子の分散状態が不均一となって熱間加工性が低下したり、熱間加工時の変形の不均一性によって、マトリックスと粒子の界面に欠陥が生じやすくなって熱間加工後のヤング率の低下が大きくなる場合がある。
【００６７】
▲１▼式及び▲２▼式で表されるＳＢとＳＣが
１００＜ＳＢ＜１０５０、
及び
５０＜ＳＣ＜５５０、
を満足する合金組成であれば、鋳造時の冷却速度が０．０５〜１℃／秒の範囲で、所望の凝固形態を安定、且つ容易に実現できる。なお、鋳造時の冷却速度が１℃／秒を超える場合には、初晶の六方晶系硼化物や立方晶系炭窒化物の粒径が２０μｍを超える場合がある。又、冷却速度が０．０５℃／秒未満では、マトリックスと分散粒子の熱収縮量の差が顕著となり、マトリックスと分散粒子の界面で欠陥が生ずる場合がある。したがって、鋳造に際しては冷却速度を０．０５〜１℃／秒として鋼塊にするのがよい。
【００６８】
なお、鋼塊は３００℃以下の温度域まで冷却するのがよい。３００℃以下の温度域まで冷却してマトリックスをマルテンサイト相に変態させると、その後の焼戻し過程で析出する炭窒化物が微細化して高強度化できるからである。更に、３００℃以下の温度域まで冷却することによって、粒界脆化を防止することができる。
【００６９】
（Ｃ−３）均質化処理条件
本発明に係る高強度高剛性鋼は、特定のオーステナイト領域で均質化処理を行うことにより、加工性及び強度を高めることができる。既に述べたように、オーステナイト域の上限（ＨＴ℃）は式中の元素記号をその元素の質量％での含有量として、前記▲４▼の実験式で表すことができるので、ＨＴ℃以下の温度で均質化処理すればよい。なお、均質化処理は、オーステナイト域上限である上記ＨＴ℃の近傍で行うことが望ましい。
【００７０】
上記均質化処理における保持時間は２〜６時間程度であればよい。
【００７１】
この均質化処理に続けて、熱間加工を行い所望の形状に仕上げてもよいし、均質化処理後に一旦冷却し、その後、均質化処理と同様にオーステナイト領域に再度加熱してから熱間加工を施して所望の形状に仕上げてもよい。
【００７２】
（Ｃ−４）熱間加工条件
熱間加工によってＦｅマトリックスの結晶粒を微細化することにより、強度及び靱性を高めることができる。
【００７３】
ここで、上記（Ｃ−３）項に述べたようにして均質化処理された鋼塊、或いは、均質化処理後に再度加熱処理された鋼塊に対しては、総減面率が５０％以上となる熱間加工を行って所望の形状に仕上げるのがよい。総減面率を５０％以上とすることで、Ｆｅマトリックスの結晶粒径は２０μｍ以下となり、強度及び靱性が向上する。なお、上記熱間加工における歪速度は、２０／秒以下が望ましい。上記熱間加工における総減面率の上限は特に規定されるものではなく、設備能力から定まる最も高い総減面率や、鋼塊と加工後の仕上げ形状との関係から定まる最も高い総減面率であってもよい。
【００７４】
ここで、上記熱間加工における総減面率とは、鋼塊を熱間加工した際の断面積の減少割合のことをいい、熱間加工前の鋼塊の断面積をＳｂ　、熱間加工後の鋼材の断面積をＳａ　として、｛（Ｓｂ−Ｓａ）／Ｓｂ　｝×１００（％）で表すことができる。
【００７５】
（Ｃ−５）焼ならし条件
一般に、金属材料に加工歪が入るとヤング率が低下する。したがって、熱間加工後、焼ならしを行い、歪を緩和させることが望ましい。
【００７６】
ここで、上記（Ｃ−４）項に述べたようにして総減面率で５０％以上の熱間加工を施された鋼片に対しては、８５０〜１０５０℃で焼ならしを行った後、３００℃以下の温度域まで冷却するのがよい。この処理により、熱間加工時に導入された加工歪が解放されるとともに、六方晶系硼化物及び立方晶系炭化物が析出し、ヤング率が向上するからである。焼ならし後は、Ｆｅマトリックス（オーステナイト）をマルテンサイト相に変態させ、高強度化を図るため上記のように３００℃以下の温度域まで冷却するのがよい。
【００７７】
なお、上記の焼ならし処理における保持時間は０．５〜２時間程度であればよい。
【００７８】
（Ｃ−６）焼戻し条件
焼戻し処理によってマルテンサイト母相にＴｉ（Ｃ、Ｎ）、Ｖ（Ｃ、Ｎ）、Ｎｂ（Ｃ、Ｎ）などの微細な炭窒化物を析出させれば、一層の高強度化を実現することができる。したがって、上記（Ｃ−５）項に述べたようにして焼ならしされた鋼片に対しては、７００℃以下で焼戻し処理をするのがよい。なお、上記焼戻し処理における保持時間は０．５〜２時間程度であればよい。
【００７９】
以下、実施例により本発明を更に詳しく説明する。
【００８０】
【実施例】
表１に示す化学組成を有する鋼を、その合金成分のうちＴｉ及びＢ以外の原料を真空中で溶解し、その後、Ｔｉ及びボロン含有合金鉄（フェロボロン）をこの順に添加して完全に溶解し、１５０ｋｇの鋼塊を得た。鋼塊のサイズは、上側直径が２３０ｍｍ、下側直径が１９０ｍｍ、高さが５００ｍｍであり、鋳造時の冷却速度は、表面部が０．５℃／秒程度、中心部が０．１℃／秒程度であった。なお、凝固した鋼塊は常温まで冷却した。
【００８１】
【表１】

【００８２】
次いで、上記のようにして得た各鋼塊を高さ２５０ｍｍの位置で２分割し、そのうち上側の鋼塊には表２に示す温度で均質化処理を行い、室温まで冷却した後、表２に示す温度で焼ならし処理及び焼戻し処理を行い、鋼塊の外周部、中心部及び、外周部と中心部の中間部の３箇所からヤング率測定用のサンプル及びミクロ試験片を採取した。なお、以下の説明では、上記の処理を施した材料を「鋳造＋熱処理」材という。
【００８３】
一方、２分割した下側の鋼塊には表２に示す温度で均質化処理を施した後、直ちに熱間鍛伸を行って直径が６０ｍｍ又は８０ｍｍの棒材を得た。この棒材を表２に示す条件で焼ならし処理、焼戻し処理を行ったのち、外周部、中心部及び、外周部と中心部の中間部の３箇所から鋳造軸に平行にヤング率測定用サンプル、引張試験片及びミクロ試験片を採取した。なお、以下の説明では、上記の処理を施した材料を「鋳造＋熱間加工＋熱処理」材という。
【００８４】
ここで、均質化処理後、焼ならし後及び焼戻し後の冷却は、いずれも大気中放冷とした。
【００８５】
【表２】

【００８６】
鏡面研磨したミクロ試料をナイタールで腐食し、走査電子顕微鏡で観察した８視野の像から、初晶型の六方晶系硼化物及び立方晶系炭窒化物（すなわち、粒径が１μｍ以上である六方晶系硼化物及び立方晶系炭窒化物）のうち粒径が最大である粒子を調査した。
【００８７】
なお、上記初晶型の六方晶系硼化物及び立方晶系炭窒化物のうちの最大粒子の粒径は、「鋳造＋熱処理」材と「鋳造＋熱間加工＋熱処理」材のいずれについても、外周部、中心部及び、外周部と中心部の中間部でほとんど差は認められず、又、「鋳造＋熱処理」材と「鋳造＋熱間加工＋熱処理」材でも差は認められなかった。
【００８８】
例として、表３及び表４に、「鋳造＋熱間加工＋熱処理」材の中心部における六方晶系硼化物及び立方晶系炭窒化物のうちの最大粒子の粒径を示す。
【００８９】
分散粒子の結晶構造の特定は次のようにして行った。すなわち、「鋳造＋熱間加工＋熱処理」材の中心部から採取したミクロ試料を電解抽出して分散粒子を抽出し、得られた抽出残渣についてＸ線回折して結晶構造を特定した。
【００９０】
分散粒子の体積分率は次のようにして求めた。
【００９１】
すなわち、　（１）「鋳造＋熱間加工＋熱処理」材の抽出残渣の定量を行い、鋼中に存在する分散粒子の総質量比を求める。　（２）次に、抽出残渣をＸ線回折し、各粒子の回折ピークを特定して各々の積分強度を求める。　（３）積分強度比から、各分散粒子の質量比を求める。　（４）予めＥＤＸ分析により求めておいた各分散粒子の組成から粒子の比重を求める。　（５）各粒子の比重を用いて、体積分率を求める。
【００９２】
表３及び表４に、上記のようにして求めた六方晶系硼化物及び立方晶系炭窒化物の体積分率を併せて示す。なお、表３及び表４においては六方晶系硼化物をＸＢ_２　、立方晶系炭窒化物をＸ（Ｃ、Ｎ）と記載し、上記ＸＢ_２　及びＸ（Ｃ、Ｎ）以外の分散粒子の体積率も併記した。
【００９３】
【表３】

【００９４】
【表４】

【００９５】
ヤング率は、厚さが１．５ｍｍ、幅が１０ｍｍで長さが６０ｍｍのヤング率測定用サンプルを用いて、通常の横共振法で共振周波数を測定することにより求めた。
【００９６】
表３及び表４に、「鋳造＋熱処理」材と「鋳造＋熱間加工＋熱処理」材のそれぞれについて、外周部、中心部及び、外周部と中心部の中間部について求めたヤング率の平均値を整理して示す。
【００９７】
ＪＩＳ　Ｚ　２２０１に記載の１４Ａ号試験片（直径が８ｍｍで評点距離が５６ｍｍ）を用いて、常温、１／秒の歪速度の条件で引張試験を行った。表３及び表４に、外周部、中心部及び、外周部と中心部の中間部について求めた０．２％耐力の平均値を整理して示す。
【００９８】
表３及び表４から、次のことが明らかである。
【００９９】
Ｃが含まれていない鋼１〜４のうち、ＴｉＢ_２　の体積分率が１０％を下回る鋼１及び鋼２の場合、初晶のＴｉＢ_２　の最大粒子の粒径はそれぞれ７．０μｍと１６．４μｍで２０μｍ以下であるがヤング率は低く、目標の２４０ＧＰａに達していない。一方、鋼３及び鋼４の場合、ＴｉＢ_２　の体積分率は１０％以上であるものの、初晶のＴｉＢ_２　の最大粒子の粒径が２０μｍを超えているので、やはりヤング率は低く、目標の２４０ＧＰａに達していない。更に、「鍛造＋熱処理」材のヤング率に比べて、「鋳造＋熱間加工＋熱処理」材のヤング率は４〜７ＧＰａ低下している。なお、上記鋼１〜４の０．２％耐力は高々２９９ＭＰａであり、強度が低い。
【０１００】
Ｂが含まれていない鋼５〜８のうち、Ｔｉ（Ｃ、Ｎ）の体積分率が１０％を下回る鋼５及び鋼６の場合、初晶のＴｉ（Ｃ、Ｎ）の最大粒子の粒径はそれぞれ４．０μｍ及び１９．０μｍと２０μｍ以下であるがヤング率は低く、目標の２４０ＧＰａに達していない。又、鋼７及び鋼８の場合、Ｔｉ（Ｃ、Ｎ）の体積分率は１０％以上であるものの、初晶のＴｉ（Ｃ、Ｎ）の最大粒子の粒径が２０μｍを超えているので、やはりヤング率は低く、目標の２４０ＧＰａに達していない。特に、熱間加工を受けた材料のヤング率、つまり「鋳造＋熱間加工＋熱処理」材のヤング率は、「鍛造＋熱処理」材のヤング率に比べて３〜７ＧＰａ低下している。なお、上記鋼５〜８の０．２％耐力は高々５５２ＭＰａである。
【０１０１】
鋼９の場合、初晶のＴｉＢ_２　とＴｉ（Ｃ、Ｎ）の最大粒子の粒径は２０μｍ以下であり、更に、ＴｉＢ_２　とＴｉ（Ｃ、Ｎ）を複合して含んでいるが、合計の体積分率が１０％に満たないので、ヤング率は低く、目標の２４０ＧＰａに達していない。
【０１０２】
これに対し、鋼１０〜２０の場合には、本発明で規定する体積割合のＸＢ_２　及びＸ（Ｃ、Ｎ）が確保され、しかも、初晶のＸＢ_２　とＸ（Ｃ、Ｎ）の最大粒子の粒径は２０μｍ以下であるため、２４０ＧＰａを超えるヤング率が得られている。
【０１０３】
具体的には、鋼１０〜１２の場合は、ＸＢ_２　とＸ（Ｃ、Ｎ）を合計の体積分率で１３〜２１．７％含有し、且つ、初晶のＸＢ_２　とＸ（Ｃ、Ｎ）の最大粒子の粒径は１１．０〜１４．８μｍと２０μｍ以下であるため、２４８〜２７５ＧＰａの高いヤング率が得られている。なお、熱間加工を受けた場合でも、ヤング率はほとんど低下してない。
【０１０４】
なお、鋼１３の場合、Ｎｂを含むので０．２％耐力が７２０ＭＰａまで増加している。
【０１０５】
鋼１４の場合は、Ｖを多量に含むことによって、（Ｔｉ、Ｖ）Ｂ_２　及び微細な（Ｔｉ、Ｖ）（Ｃ、Ｎ）が高密度に分散しており、０．２％耐力で９９８ＭＰａという極めて高い強度が得られている。
【０１０６】
鋼１５の場合は、Ｎｂを多量に含むことによって、（Ｔｉ、Ｎｂ）Ｂ_２　及び微細な（Ｔｉ、Ｎｂ）（Ｃ、Ｎ）が高密度に分散しており、０．２％耐力で８８６ＭＰａという高い強度が得られている。
【０１０７】
鋼１６の場合は、Ｖ及びＮｂを複合して含むので、０．２％耐力で９７０ＭＰａという高い強度が得られている。
【０１０８】
鋼１７の場合は、Ｎｄを含むことにより不純物元素が固定され、分散粒子とＦｅマトリックスとの密着性が向上するので、熱間加工後のヤング率が向上している。
【０１０９】
鋼１８の場合は、Ｖを多量に含むことによって、高いヤング率と０．２％耐力で１０１１ＭＰａという極めて高い強度が得られている。
【０１１０】
鋼１９の場合は、ＣａとＮｄを複合して含むので、熱間加工後にヤング率が向上している。
【０１１１】
鋼２０の場合は、Ｖを多量に含む硬質の鋼であるが、Ｍｇを含んでいるので、熱間加工による棒材の仕上げ直径をより小さくした場合、すなわち総減面率が大きい場合でもヤング率が低下していない。
【０１１２】
なお、鋼２１〜２３の場合、マトリックス中に分散する粒子の量は多いが、高ヤング率が得られない。
【０１１３】
すなわち、鋼２１の場合は、ＴｉＢ_２　とＴｉ（Ｃ、Ｎ）の体積分率は２５．９％であるものの、初晶のＴｉＢ_２　とＴｉ（Ｃ、Ｎ）の最大粒子の粒径が２０μｍを超えて２４．２μｍであるので、ヤング率は低く、目標の２４０ＧＰａに達していない。
【０１１４】
鋼２２の場合は、ＴｉＢ_２　とＴｉ（Ｃ、Ｎ）の合計体積分率が５．２％と低く、しかもＦｅマトリックスに分散している粒子の１５％がＦｅ_３Ｃ（セメンタイト）であり、Ｆｅ_３Ｃのヤング率はＦｅマトリックスより低いため、ヤング率は極めて低い。
【０１１５】
鋼２３の場合は、ＴｉＢ_２　とＴｉ（Ｃ、Ｎ）の合計体積分率が２．０％未満と低く、しかも高ヤング率化に寄与しないＦｅ_３Ｃを９％、斜方晶系硼化物を１５％含むため、ヤング率は極めて低い。
【０１１６】
【発明の効果】
本発明の高強度高剛性鋼は、溶鋼を鋳造して製造しても、晶出粒子の粗大化が抑制され、熱間加工による成形が可能であって、しかも、２４０ＧＰａ以上のヤング率を有するので、例えば、自動車エンジン用各種部品や、ボルト、各種シャフト類といった大型の構造部材にも利用することができる。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a high-strength, high-rigidity steel suitable for structural members that require rigidity, such as materials for automobiles, materials for robots, and materials for sports equipment, and a method for manufacturing the same.
[0002]
[Prior art]
For example, in the case of materials for automobiles, there is an increasing demand for high-strength and high-rigidity materials for the purpose of reducing the weight for improving fuel consumption and improving riding comfort.
[0003]
As an index representing rigidity, Young's modulus can be cited. Conventionally, in the case of an Fe-based alloy (hereinafter also referred to as steel), alloying, that is, isotropic Young's modulus is increased by adding an alloy element. In addition, the Young's modulus in a specific direction has been increased by developing a texture.
[0004]
However, the elements that contribute to the improvement of the Young's modulus of steel by alloying are limited to Cr, Co, and Re, and even when these elements are contained in the steel, the Young's modulus is improved only by a few percent at most. On the other hand, when a texture is used to improve the Young's modulus of steel, the Young's modulus has a large anisotropy, and therefore there is a limit to application to structural members.
[0005]
In recent years, composite materials having a high Young's modulus in which particles having a high Young's modulus are dispersed in a matrix have been studied.
[0006]
In general, the Young's modulus of a composite material is as described in the “elastic modulus of metal material” edited by the Japan Society of Mechanical Engineers (issue date: October 31, 1980, publisher: Japan Society of Mechanical Engineers). It is determined by the volume fraction of the dispersed particles. The larger the Young's modulus of the dispersed particles and the greater the content thereof, the greater the improvement. That is, according to the composite law of Young's modulus, when the matrix is α-Fe (ferrite), for example, about 25 vol% SiC, about 30% vol% Al ₂ O ₃ About 15 vol% TiB ₂ By dispersing any of about 20 vol% TiC in the matrix, the Young's modulus becomes 250 GPa, which is improved by about 20% from the normal value of 210 GPa.
[0007]
However, in order to disperse particles having a high Young's modulus in the above ratio in the ferrite matrix, it is necessary to add a large amount of alloying elements. Therefore, the high rigidity is obtained through normal melting and hot working processes. It is difficult to produce steel.
[0008]
For this reason, in order to produce high-rigidity steel by a melting method, for example, as in the technique disclosed in Japanese Patent Application Laid-Open No. 10-68048, a temperature at which elements constituting a compound are completely dissolved in iron or an iron alloy The compound is crystallized or precipitated at the time of heating, cooling and solidification, or alloy raw materials are blended at a predetermined ratio as in the technique disclosed in Japanese Patent Application Laid-Open No. 2001-59146, and in vacuum or not. After completely melting in an active gas atmosphere, it is necessary to cast into a mold or a ceramic mold.
[0009]
However, even in the case of steel manufactured in this way, if the primary crystal particles crystallized from the molten steel agglomerate and coarsen, the adhesion at the interface between the dispersed particles and the matrix decreases, and there are defects such as cracks and voids at the interface. If such defects occur in the material, a large decrease in Young's modulus is inevitable. Further, when deformation such as hot working is applied to the material in which the defect has occurred, the defect increases, and thus the Young's modulus is further decreased. For this reason, in the case of molten steel containing 10% by volume fraction of high-rigidity particles, the composite law of Young's modulus does not hold, and particularly in steel that has undergone processes such as hot working, a desired value such as 240 GPa is desired. However, there is a problem that ductility and toughness become extremely low due to the influence of coarse particles.
[0010]
Therefore, for the production of high-rigidity steel, for example, as disclosed in Japanese Patent Application Laid-Open No. 5-239504, a method using a mechanical alloying method has been mainstream. However, the mechanical alloying method has a problem that the cost increases and large parts cannot be manufactured.
[0011]
[Problems to be solved by the invention]
The present invention has been made in view of the above-mentioned present situation, and its purpose is to form by hot working even for steel containing 10% or more of high-rigidity particles produced by a melting method in volume fraction, Moreover, it is to provide a high-strength and high-rigidity steel having a Young's modulus of 240 GPa or more and suitable as a structural member such as an automobile material, a robot material, and a sports equipment material, and a method for producing the same.
[0012]
[Means for Solving the Problems]
The gist of the present invention resides in the high-strength high-rigidity steel shown in the following (1) to (3) and the high-strength high-rigidity steel shown in (4).
[0013]
(1) A high-strength, high-rigidity steel manufactured by casting molten steel, in which a hexagonal boride and cubic carbonitride are dispersed in a total volume of 10 to 35% in the Fe-based alloy matrix. And high-strength and high-stiffness steel, wherein the hexagonal boride and cubic carbonitride have a particle size of 20 μm or less.
[0014]
(2) Hexagonal boride is TiB ₂ , VB ₂ , NbB ₂ And one or more of these composite borides, and cubic carbonitrides are Ti (C, N), V (C, N), Nb (C, N), Ta (C, N) Hf (C, N) and one or more of these composite carbonitrides, the high-strength and high-stiffness steel according to (1) above.
[0015]
(3) By mass%, C: 0.2 to 1.7%, Si: 0.05 to 0.5%, Mn: 0.2 to 1.5%, Ti: 1 to 12%, Mo: 0 0.1 to 2%, B: 0.5 to 3.2%, N: 0.001 to 0.01%, V: 0 to 12%, Nb: 0 to 12%, Ta: 0 to 1%, Hf : 0 to 1%, Ca: 0 to 0.007%, Mg: 0 to 0.007%, and Nd: 0 to 0.007%, with the balance being Fe and impurities. The high-strength, high-rigidity steel as described in (2) above, wherein SB represented is more than 100 and less than 1050, and SC represented by the following formula (2) is more than 50 and less than 550.
[0016]

However, the element symbol in each formula represents the content in mass% of the element.
[0017]
(4) By mass%, C: 0.2 to 1.7%, Si: 0.05 to 0.5%, Mn: 0.2 to 1.5%, Ti: 1 to 12%, Mo: 0 0.1 to 2%, B: 0.5 to 3.2%, N: 0.001 to 0.01%, V: 0 to 12%, Nb: 0 to 12%, Ta: 0 to 1%, Hf : 0 to 1%, Ca: 0 to 0.007%, Mg: 0 to 0.007% and Nd: 0 to 0.007%, with the balance being Fe and impurities, Casting process of molten steel in which a compound of each element in steel in which SB represented is more than 100 and less than 1050 and SC represented by the formula (2) is more than 50 and less than 550 is dissolved in the steel First, hexagonal borides or cubic carbonitrides are crystallized, and then the three phases of hexagonal borides, cubic carbonitrides and Fe matrix phase are crystallized from the residual molten steel. The steel ingot is cooled to a temperature range of 300 ° C. or lower, further homogenized at a temperature of HT (° C.) or less expressed by the following formula (4), and then HT (° C.). Hot working is performed so that the total area reduction rate is 50% or more at the following temperatures, followed by further normalizing at 850 to 1050 ° C., cooling to 300 ° C. or lower, and then tempering at 700 ° C. or lower. A method for producing high-strength, high-rigidity steel.
[0018]

However, the element symbol in the formula (4) represents the content in mass% of the element.
[0019]
The “particle diameter” in the present invention specifically refers to a value defined by ½ of the sum of the short diameter and long diameter of each particle.
[0020]
The hexagonal boride and cubic carbonitride specified in the present invention are obtained by, for example, extracting a crystallized product or precipitate in steel by a usual method such as electrolytic extraction, and then obtaining the residue by X-ray. This can be confirmed by diffraction.
[0021]
TiB ₂ , VB ₂ And NbB ₂ The compound boride of “XB” is “XB” in which any two or more selected from Ti, V and Nb are used. ₂ , Specifically, for example, (Ti, V) B ₂ , (V, Nb) B ₂ Ya (Ti, V, Nb) B ₂ And so on. Also, Ti (C, N), V (C, N), Nb (C, N), Ta (C, N) and Hf (C, N) composite carbonitrides are “X” for Ti, Carbonitride represented by “X (C, N)” as any two or more selected from V, Nb, Ta and Hf, specifically, for example, (Ti, V) (C, N) , (V, Nb, Hf) (C, N), (Ti, V, Nb, Hf) (C, N), (Ti, V, Nb, Hf, Ta) (C, N), and the like.
[0022]
The hexagonal borides and cubic carbonitrides described above are classified into the following three types (a) to (c) depending on the form of formation.
[0023]
(B) Primary crystal hexagonal boride and cubic carbonitride:
It is first crystallized from molten steel in the cooling process of casting, and its particle size is 1 μm or more.
[0024]
(B) Eutectic hexagonal boride and cubic carbonitride:
After crystallization of the primary crystal type boride or carbonitride of (b) above, the boride, carbide and Fe matrix are crystallized almost simultaneously from the residual molten steel, and the particle size is less than 1 μm. It is.
[0025]
(C) Precipitated hexagonal boride and cubic carbonitride:
It is formed in a cooling process after the steel is completely solidified, or in a process such as a heat treatment performed after the solidification, and its particle size is 0.1 μm or less.
[0026]
Here, the crystallized hexagonal boride and cubic carbonitride of (b) and (b) can be observed using an optical microscope or a scanning electron microscope, and (c) precipitation type These hexagonal borides and cubic carbonitrides can be observed, for example, by using a transmission electron microscope having an acceleration voltage of 100 to 200 kV.
[0027]
The present inventors stably and finely crystallize or precipitate even when manufactured by a melting method, and ensure a high Young's modulus not only as an as-cast material but also by hot working after casting. In order to find the dispersed particles that can be used, in addition to the chemical composition of the steel, various conditions such as casting conditions, heat treatment conditions, and hot working were changed. As a result, the following findings (a) to (i) were obtained.
[0028]
(A) Among the compounds produced during casting of molten steel, hexagonal borides and cubic carbonitrides have a high Young's modulus. That is, in the case of molten steel, in compounds other than the above hexagonal borides and cubic carbonitrides, Fe is distributed in the compound, and the Young's modulus of the compound itself decreases. On the other hand, even in the case of molten steel, Fe is hardly distributed in the hexagonal boride and the cubic carbonitride, so that a high Young's modulus is ensured. Therefore, hexagonal borides and cubic carbonitrides are effective for increasing the Young's modulus of molten steel.
[0029]
(B) Among hexagonal borides, “X” is one or more selected from Ti, V, and Nb, and “XB ₂ The hexagonal boride represented by the formula has the highest Young's modulus and is suitable for increasing the Young's modulus of molten steel.
[0030]
(C) Among cubic carbonitrides, “X” is one or more selected from Ti, V, Nb, Ta and Hf, and the cubic system represented by “X (C, N)” Carbonitride has the highest Young's modulus and is effective for increasing the Young's modulus of molten steel.
[0031]
(D) The cubic carbonitride obtained by precipitation also has a strong precipitation strengthening action.
[0032]
(E) In order to achieve high Young's modulus of molten steel with either hexagonal boride or cubic carbonitride, B in the former case, C in the latter case, It is necessary to add together with a large amount of alloying elements. In these cases, in the cooling process when casting molten steel, coarse hexagonal boride or cubic carbonitriding with a particle size exceeding 20 μm in the molten steel. The product crystallizes, and these agglomerate and coarsen in the molten steel. Such a coarse hexagonal boride or cubic carbonitride having a particle size exceeding 20 μm has low adhesion to the interface of the Fe matrix, so that the Young's modulus of the steel is lowered. Further, when molding is performed by applying plastic deformation such as hot working after casting, defects occur between the hexagonal boride or cubic carbonitride and the Fe matrix, and the Young's modulus further decreases. For this reason, in the above case, when the volume fraction of particles exceeds 10% in an as-cast material or hot-worked material, the composite law of Young's modulus does not hold and, for example, a high Young's modulus of 240 GPa is realized Is difficult.
[0033]
(F) When either one or both of hexagonal boride and cubic carbonitride crystallizes from molten steel containing both C and B, C is also added to the hexagonal boride, Since B lowers the interfacial energy for cubic carbonitrides, aggregation and coarsening of primary hexagonal borides and / or primary cubic carbonitrides are suppressed. Furthermore, after primary hexagonal borides and / or cubic carbonitrides crystallize, hexagonal borides, cubic carbonitrides and Fe matrix phases are crystallized almost simultaneously from low-temperature residual molten steel. By crystallization (eutectic), the particle size of the crystallization particles becomes 20 μm or less. As a result, the Young's modulus does not decrease not only in the as-cast material but also when hot working is subsequently applied, and the Young's modulus composite law is established even when the volume fraction of the particles is in the range of 10 to 35%.
[0034]
(G) In order to realize the solidification form described in (f) above under normal cooling conditions during casting, for example, within a cooling rate range of 0.05 to 1 ° C./second, hexagonal boride and molten steel The temperature range (solid-liquid coexistence region) in which the coexisting with each other and the temperature range (solid-liquid coexistence region) in which the cubic carbonitride and molten steel coexist may be set as follows.
[0035]
-100 ° C <coexisting temperature range of molten steel and hexagonal boride <1050 ° C,
-50 ° C <coexisting temperature range of molten steel and cubic carbonitride <550 ° C.
[0036]
Then, about the said temperature range, the differential thermal analysis and the frozen structure | tissue from each temperature range were observed, and the influence of the component element was investigated. As a result, it became clear that the coexisting temperature range of the molten steel and hexagonal boride and the coexisting temperature range of the molten steel and cubic carbonitride can be expressed by the following SB index and SC index, respectively.
[0037]
That is, in order to realize the solidification form described in the above (f), the SB represented by the following formula (1) exceeds 100 when the element symbol in each formula is the content in mass% of the element. The SC represented by the following formula (2) may be more than 50 and less than 550.
[0038]

[0039]
(H) In the case of having a specific chemical composition and satisfying the condition (g), a decrease in Young's modulus can be minimized by performing hot working in the austenite stable region. The upper limit HT (° C.) of the austenite region can be expressed by the empirical formula (4) above, where the element symbol in the formula is the content in mass% of the element. What is necessary is just to process.
[0040]
(I) By carrying out a tempering treatment at a temperature of 700 ° C. or less, fine carbides are precipitated and the strength of the steel can be increased.
[0041]
The present invention has been completed based on the above findings.
[0042]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, each requirement of the present invention will be described in detail.
[0043]
(A) Particles in the matrix of the Fe-based alloy
In order to increase the Young's modulus and ensure that the steel has a Young's modulus of 240 GPa or higher, hexagonal borides and cubic carbonitrides must be dispersed in a Fe-based alloy matrix by a total volume of 10% or more. I must. On the other hand, if the total volume fraction of hexagonal borides and cubic carbonitrides dispersed in the Fe-based alloy matrix exceeds 35%, the Young's modulus decreases and cracks or porosity in the steel ingot Occurs, and subsequent processing becomes difficult. Accordingly, in the present invention, hexagonal borides and cubic carbonitrides are dispersed in a total volume of 10 to 35% in the matrix of the Fe-based alloy.
[0044]
The hexagonal borides and cubic carbonitrides specified in the present invention include the above-described primary hexagonal borides and cubic carbonitrides having a particle size of 1 μm or more, and a particle size of 1 μm. Both eutectic hexagonal boride and cubic carbonitride having a particle size of less than 0.1 μm and precipitation type hexagonal boride and cubic carbonitride having a particle size of 0.1 μm or less are included.
[0045]
Here, the composition, crystal structure and volume ratio of a certain phase are determined by electrolytic extraction of particles other than the Fe matrix in steel, quantitative chemical analysis of the obtained residue, X-ray diffraction intensity measurement, and EDX analysis of individual particles. And by measuring the density of the steel.
[0046]
The aforementioned crystallized and precipitated hexagonal borides and cubic carbonitrides should be observed using an optical microscope, a scanning electron microscope, and a transmission electron microscope having an acceleration voltage of 100 to 200 kV, for example. Therefore, the image obtained by observation is image-analyzed to measure the short diameter and the long diameter, and the particle diameter of each particle can be obtained from 1/2 of the sum. With respect to hexagonal borides and cubic carbonitrides, when the particle size obtained as described above exceeds 20 μm, it is difficult to stably secure a desired high Young's modulus of 240 GPa or more. . Therefore, the particle size of hexagonal boride and cubic carbonitride dispersed in a total volume of 10 to 35% in the matrix of the Fe-based alloy was set to 20 μm or less. The lower limit value of the particle size of the hexagonal boride and cubic carbonitride is not necessarily specified, but the lower limit may be about 5 nm that can be observed with a transmission electron microscope having an acceleration voltage of 100 to 200 kV. .
[0047]
The ratio of hexagonal boride and cubic carbonitride dispersed in a total volume of 10 to 35% in the matrix of the Fe-based alloy is preferably 5: 1 to 1: 1. By satisfying this condition, an extremely healthy steel ingot can be obtained.
[0048]
Hexagonal borides as particles dispersed in a Fe-based alloy matrix with a volume fraction of 10 to 35% together with cubic carbonitrides are TiB ₂ , VB ₂ , NbB ₂ And one or more of these composite borides. TiB ₂ , VB ₂ And NbB ₂ The compound boride of “XB” is “XB” in which any two or more selected from Ti, V and Nb are used. ₂ , Specifically, for example, (Ti, V) B ₂ , (V, Nb) B ₂ Ya (Ti, V, Nb) B ₂ It is as already mentioned that it points out.
[0049]
Further, cubic carbonitrides as particles dispersed in a Fe-based alloy matrix together with hexagonal borides in a volume fraction of 10 to 35% are Ti (C, N), V (C, N). Nb (C, N), Ta (C, N), Hf (C, N) and one or more of these composite carbonitrides are preferable. As already described, the composite carbonitride of Ti (C, N), V (C, N), Nb (C, N), Ta (C, N) or Hf (C, N) is “X” is any two or more selected from Ti, V, Nb, Ta and Ha, and a carbonitride represented by “X (C, N)”, specifically, for example, (Ti, V ) (C, N), (V, Nb, Hf) (C, N), (Ti, V, Nb, Hf) (C, N) and (Ti, V, Nb, Hf, Ta) (C, N ) Etc.
[0050]
(B) Chemical composition of steel
Next, a preferable chemical composition of the high strength and high rigidity steel according to the present invention will be described. In addition, "%" display of the content of each element means "mass%".
[0051]
C: 0.2-1.7%
C is an element effective for forming cubic carbonitride and increasing the Young's modulus and strength of steel. Furthermore, it has the effect | action which suppresses the aggregation coarsening of the boride crystallized in molten steel. However, when the content is less than 0.2%, the crystallization amount of cubic carbonitride is small, and the desired Young's modulus of 240 GPa or more may not be obtained. On the other hand, if the content exceeds 1.7%, the crystallized carbonitride tends to be coarsened, and the desired Young's modulus cannot be obtained, and the castability and hot workability are adversely affected. There is. Therefore, the content of C is preferably 0.2 to 1.7%. The C content is more preferably 0.3 to 1.2%, and very preferably 0.4 to 0.8%.
[0052]
Si: 0.05-0.5%
Si is an element effective for deoxidation, but if its content is less than 0.05%, it may be difficult to obtain the effect. On the other hand, if the steel contains a large amount of Ti, V, Nb, etc., intermetallic compounds such as Laves phase are likely to be formed at the grain boundaries, and particularly if the Si content exceeds 0.5% The compound may become coarse. An intermetallic compound such as a Laves phase causes a large decrease in hot workability, and the Young's modulus of steel decreases because the Laves phase itself is a low Young's modulus particle. Therefore, the Si content is preferably 0.05 to 0.5%. The upper limit of the Si content is more preferably 0.3%, and even more preferably 0.2%.
[0053]
Mn: 0.2 to 1.5%
The high-strength and high-rigidity steel according to the present invention is improved in workability and strength by performing a homogenization treatment in a specific austenite region. For this reason, it is preferable to expand the austenite region using Mn, which is an austenite stabilizing element. Mn also has the effect of fixing S, which is an impurity of steel, to improve hot workability. However, if the Mn content is less than 0.2%, the above effect may not be obtained. On the other hand, if the content exceeds 1.5%, M decreases the Young's modulus. ₂₃ C ₆ Mold carbides are likely to form. Therefore, the Mn content is preferably 0.2 to 1.5%. The Mn content is more preferably 0.8 to 1.2%.
[0054]
Ti: 1 to 12%
Ti is TiB ₂ , (Ti, V) B ₂ This is an element effective for increasing the Young's modulus of steel by forming hexagonal borides such as Ti and cubic carbonitrides such as Ti (C, N). Furthermore, it is effective for increasing the strength of steel. However, when the content is less than 1%, the above effect may not be obtained. On the other hand, if the content exceeds 12%, hexagonal borides and cubic carbonitrides may become coarse and the Young's modulus may decrease. Therefore, the Ti content is preferably 1 to 12%. The Ti content is more preferably 3 to 10%, even more preferably 5 to 9%, and particularly preferably 7 to 8%.
[0055]
Mo: 0.1 to 2%
When Mo is present in a solid solution state in the matrix, it has an effect of strengthening grain boundaries and remarkably improving hot workability. Furthermore, it also has the effect of increasing the strength of the steel. However, when the content is less than 0.1%, the above effect may not be obtained. On the other hand, even if Mo is contained in excess of 2%, the effect may be saturated and the cost will increase. Therefore, the Mo content is preferably 0.1 to 2%. A more desirable range of the Mo content is 0.2 to 1%, and 0.3 to 0.7% is extremely preferable.
[0056]
B: 0.5-3.2%
B is an element effective for forming a hexagonal boride and increasing the Young's modulus of the steel. However, when the content is less than 0.5%, the amount of hexagonal boride formed is small, and the desired Young's modulus of 240 GPa or more may not be obtained. On the other hand, if the content exceeds 3.2%, the borides that crystallize out become coarse, the Young's modulus decreases, and hot workability may be adversely affected. Therefore, the content of B is preferably 0.5 to 3.2%. The content of B is more preferably 1.0 to 3.0%, even more preferably 1.8 to 2.7%, and extremely preferably 2.0 to 2.4%.
[0057]
N: 0.001 to 0.01%
N is an element effective for forming cubic carbonitride and increasing the Young's modulus of steel. However, when the content is less than 0.001%, the above effect may not be obtained. On the other hand, when the content exceeds 0.01%, the hot workability may be deteriorated. Therefore, the N content is preferably 0.001 to 0.01%. A more desirable range of the N content is 0.003 to 0.007%.
[0058]
The high-strength and high-rigidity steel according to the present invention may be one or more selected from Nb, V, Hf and Ta, or / and Ca, Mg and One or more selected from Nd may be selectively contained.
[0059]
V: 0-12%
If V is added, VB ₂ , (Ti, V) B ₂ And hexagonal borides such as V (C, N), (Ti, V) (C, N), etc. are formed to increase the Young's modulus of the steel. It also has the effect of increasing the strength of the steel. In order to surely obtain these effects, it is preferable that V has a content of 0.5% or more. However, if its content exceeds 12%, the workability may be significantly reduced. Therefore, the content of V is preferably 0 to 12%.
[0060]
Nb: 0 to 12%
If Nb is added, NbB ₂ , (Ti, Nb) B ₂ And hexagonal borides such as Nb (C, N), (Ti, Nb) (C, N), and the like, and have the effect of increasing the Young's modulus of the steel. It also has the effect of increasing hot workability and the strength of steel. In order to reliably obtain these effects, it is preferable that Nb has a content of 0.05% or more. However, if its content exceeds 12%, the workability may be significantly reduced. Therefore, the Nb content is preferably 0 to 12%.
[0061]
Ta: 0 to 1%
When Ta is added, cubic carbonitrides such as Ta (C, N) and (Ti, Ta) (C, N) are formed to increase the Young's modulus of the steel. The cubic carbonitride has the effect of refining the solidification structure and refining the crystal grains, and through this action, reduces grain boundary segregation, and thus has the effect of improving hot workability. Furthermore, it is effective for increasing the strength of steel. In order to reliably obtain these effects, Ta is preferably contained in a content of 0.05% or more. However, when the content exceeds 1%, workability may be deteriorated. Therefore, the content of Ta is preferably 0 to 1%.
[0062]
Hf: 0 to 1%
If added, Hf forms cubic carbonitrides such as Hf (C, N), (Ti, Hf) (C, N), and has the effect of increasing the Young's modulus of the steel. The cubic carbonitride has the effect of refining the solidification structure and refining the crystal grains, and through this action, reduces grain boundary segregation, and thus has the effect of improving hot workability. Furthermore, it is effective for increasing the strength of steel. In order to reliably obtain these effects, it is preferable that Hf has a content of 0.2% or more. However, when the content exceeds 1%, workability may be deteriorated. Therefore, the content of Hf is preferably 0 to 1%.
Ca: 0 to 0.007%
When Ca is added, it has an effect of fixing S and O (oxygen), which are impurities of steel, and improving castability and hot workability. In order to reliably obtain this effect, the Ca content is preferably 0.002% or more. However, if the content exceeds 0.007%, the inclusions may become coarse and the hot workability may be reduced. Therefore, the Ca content is preferably 0 to 0.007%.
Mg: 0 to 0.007%
If Mg is added, it has the effect of fixing S and O (oxygen), which are impurities of steel, and improving castability and hot workability. In order to reliably obtain this effect, the Mg content is preferably 0.002% or more. However, if the content exceeds 0.007%, the inclusions may become coarse and the hot workability may be reduced. Therefore, the content of Mg is preferably 0 to 0.007%.
Nd: 0 to 0.007%
When added, Nd fixes S and O (oxygen), which are impurities of steel, and has an effect of improving castability and hot workability. In order to reliably obtain this effect, it is preferable that Nd has a content of 0.002% or more. However, if the content exceeds 0.007%, the inclusions may become coarse and the hot workability may be reduced. Therefore, the Nd content is preferably 0 to 0.007%.
(C) Steel production method
Next, an example of manufacturing conditions for the high-strength and high-rigidity steel according to the present invention will be described.
[0063]
(C-1) Melting conditions
In order to obtain the structure of the high-strength and high-rigidity steel according to the present invention, in addition to the chemical composition described in the above section (B), the formula (1) relating to the coexisting temperature range of the molten steel and the hexagonal boride is used. SC represented by the above formula (2) relating to the coexisting temperature range of SB and molten steel and cubic carbonitride,
100 <SB <1050,
50 <SC <550,
It is desirable to select an alloy composition that satisfies the above conditions, completely dissolve the compound of each element in steel, and then perform casting so as to obtain an appropriate solidification form.
[0064]
When the value of SB is 1050 or more, the particle size of the primary hexagonal boride exceeds 20 μm, and the Young's modulus of the as-cast material or hot-worked material may be decreased. On the other hand, when the value of SB is 100 or less, the amount of hexagonal boride crystallized from the molten steel becomes insufficient, and the desired Young's modulus of 240 GPa or more may not be obtained.
[0065]
On the other hand, when the SC value is 550 or more, the grain size of the primary cubic carbonitride exceeds 20 μm, which may lead to a decrease in the Young's modulus of the as-cast material or hot-worked material. On the other hand, when the SC value is 50 or less, the amount of cubic carbonitride crystallized from the molten steel becomes insufficient, and the desired Young's modulus of 240 GPa or more may not be obtained.
[0066]
(C-2) Casting conditions
In addition to the chemical composition described in the item (B), SB and SC represented by the formulas (1) and (2)
100 <SB <1050,
as well as
50 <SC <550,
In the cooling process at the time of casting after the compound of each element in the steel satisfying the requirements is completely dissolved in the steel, first, one or both of hexagonal boride and cubic carbonitride are crystallized. Next, it is desirable to crystallize (eutectic) the three phases of the hexagonal boride, the cubic carbonitride, and the Fe matrix phase almost simultaneously from the residual molten steel. When the primary crystal is an Fe matrix such as δ ferrite, orthorhombic boride having a low Young's modulus and carbide having a high Fe content are stably crystallized or precipitated, and hexagonal boron as defined in the present invention. In some cases, the desired Young's modulus of 240 GPa or more cannot be obtained because the chemical compound or cubic carbonitride is not sufficiently crystallized. In addition, if the three phases of hexagonal boride, cubic carbonitride, and Fe matrix phase are not subjected to almost simultaneous crystallization (eutectic) process, the dispersed state of the particles becomes non-uniform and hot working is performed. There is a case where defects are easily generated at the interface between the matrix and the particles due to non-uniformity of deformation during hot working and the Young's modulus after hot working is greatly reduced.
[0067]
SB and SC represented by the formulas (1) and (2)
100 <SB <1050,
as well as
50 <SC <550,
If the alloy composition satisfies the above, a desired solidification form can be realized stably and easily when the cooling rate during casting is in the range of 0.05 to 1 ° C./second. When the cooling rate at the time of casting exceeds 1 ° C./second, the grain size of the primary crystal hexagonal boride or cubic carbonitride may exceed 20 μm. On the other hand, when the cooling rate is less than 0.05 ° C./second, the difference in thermal shrinkage between the matrix and the dispersed particles becomes significant, and defects may occur at the interface between the matrix and the dispersed particles. Therefore, it is preferable to form a steel ingot at a cooling rate of 0.05 to 1 ° C./second during casting.
[0068]
In addition, it is good to cool a steel ingot to the temperature range below 300 degreeC. This is because if the matrix is transformed to the martensite phase by cooling to a temperature range of 300 ° C. or lower, the carbonitrides precipitated in the subsequent tempering process can be refined and increased in strength. Furthermore, grain boundary embrittlement can be prevented by cooling to a temperature range of 300 ° C. or lower.
[0069]
(C-3) Homogenization treatment conditions
The high-strength, high-rigidity steel according to the present invention can improve workability and strength by performing a homogenization treatment in a specific austenite region. As already stated, the upper limit (HT ° C.) of the austenite region can be expressed by the empirical formula (4) above, where the element symbol in the formula is the content in mass% of the element. What is necessary is just to homogenize at temperature. The homogenization treatment is desirably performed in the vicinity of the above-mentioned HT ° C., which is the upper limit of the austenite region.
[0070]
The holding time in the said homogenization process should just be about 2 to 6 hours.
[0071]
Subsequent to this homogenization treatment, hot processing may be performed to finish the desired shape. After the homogenization treatment, the workpiece is once cooled and then heated again to the austenite region in the same manner as the homogenization treatment and then hot working. To give a desired shape.
[0072]
(C-4) Hot working conditions
Strength and toughness can be increased by refining Fe matrix crystal grains by hot working.
[0073]
Here, for the steel ingot that has been homogenized as described in the above section (C-3), or the steel ingot that has been heat-treated again after the homogenization, the total area reduction ratio is 50% or more. It is good to finish in a desired shape by performing hot working. By setting the total area reduction rate to 50% or more, the crystal grain size of the Fe matrix becomes 20 μm or less, and the strength and toughness are improved. The strain rate in the hot working is preferably 20 / second or less. The upper limit of the total area reduction in the above hot working is not particularly specified, the highest total area reduction determined from the equipment capacity and the highest total area reduction determined from the relationship between the steel ingot and the finished shape after processing It may be a rate.
[0074]
Here, the total area reduction ratio in the hot working refers to a reduction ratio of the cross-sectional area when the steel ingot is hot worked, and the cross-sectional area of the steel ingot before hot working is Sb, The cross-sectional area of the subsequent steel material can be represented by {(Sb-Sa) / Sb} × 100 (%), where Sa.
[0075]
(C-5) Normalizing conditions
Generally, when processing strain enters a metal material, Young's modulus decreases. Therefore, it is desirable to normalize and reduce strain after hot working.
[0076]
Here, normalization was performed at 850 to 1050 ° C. on the steel pieces subjected to hot working with a total area reduction rate of 50% or more as described in the above section (C-4). Then, it is good to cool to the temperature range below 300 degreeC. This is because the processing strain introduced during hot working is released by this treatment, and hexagonal borides and cubic carbides are precipitated to improve the Young's modulus. After normalization, the Fe matrix (austenite) is transformed into a martensite phase and cooled to a temperature range of 300 ° C. or lower as described above in order to increase the strength.
[0077]
In addition, the holding time in said normalization process should just be about 0.5 to 2 hours.
[0078]
(C-6) Tempering conditions
If fine carbonitrides such as Ti (C, N), V (C, N), and Nb (C, N) are precipitated in the martensite matrix by tempering, higher strength can be achieved. Can do. Therefore, it is preferable to temper the steel pieces normalized as described in the above section (C-5) at 700 ° C. or lower. The holding time in the tempering process may be about 0.5 to 2 hours.
[0079]
Hereinafter, the present invention will be described in more detail with reference to examples.
[0080]
【Example】
Steels having the chemical composition shown in Table 1 were melted in vacuum with raw materials other than Ti and B among the alloy components, and then Ti and boron-containing alloy iron (ferroboron) were added in this order and completely dissolved. A steel ingot of 150 kg was obtained. The steel ingot has an upper diameter of 230 mm, a lower diameter of 190 mm, and a height of 500 mm. The cooling rate during casting is about 0.5 ° C./second for the surface portion and 0.1 ° C./second for the center portion. It was about a second. The solidified steel ingot was cooled to room temperature.
[0081]
[Table 1]

[0082]
Next, each steel ingot obtained as described above was divided into two at a height of 250 mm, and the upper steel ingot was homogenized at the temperature shown in Table 2 and cooled to room temperature. A normalizing treatment and a tempering treatment were performed at the temperature shown in FIG. 5, and samples for measuring Young's modulus and micro test pieces were collected from three locations, the outer peripheral portion, the central portion, and the intermediate portion between the outer peripheral portion and the central portion. In the following description, the material subjected to the above treatment is referred to as a “casting + heat treatment” material.
[0083]
On the other hand, the lower steel ingot divided into two parts was homogenized at the temperature shown in Table 2, and immediately subjected to hot forging to obtain a bar having a diameter of 60 mm or 80 mm. After normalizing and tempering the bar under the conditions shown in Table 2, the Young's modulus is measured in parallel to the casting axis from the outer periphery, the center, and the center between the outer periphery and the center. Samples, tensile specimens and micro specimens were taken. In the following description, the material subjected to the above treatment is referred to as “casting + hot working + heat treatment” material.
[0084]
Here, after the homogenization treatment, cooling after normalization and tempering were all allowed to cool in the atmosphere.
[0085]
[Table 2]

[0086]
A mirror-polished micro sample was corroded with nital, and from an eight-field image observed with a scanning electron microscope, primary crystal hexagonal boride and cubic carbonitride (that is, hexagonal with a particle size of 1 μm or more) The particles having the largest particle size among the crystal borides and the cubic carbonitrides were investigated.
[0087]
The maximum particle size of the primary crystal type hexagonal boride and cubic carbonitride is the same for both “casting + heat treatment” material and “casting + hot working + heat treatment” material. There was almost no difference between the outer peripheral part, the central part, and the intermediate part between the outer peripheral part and the central part, and no difference was found between the “casting + heat treatment” material and the “casting + hot working + heat treatment” material. .
[0088]
As an example, Tables 3 and 4 show the particle sizes of the largest particles of hexagonal boride and cubic carbonitride in the center of the “casting + hot working + heat treatment” material.
[0089]
The crystal structure of the dispersed particles was specified as follows. That is, a micro sample collected from the center of the “casting + hot working + heat treatment” material was subjected to electrolytic extraction to extract dispersed particles, and the obtained extraction residue was subjected to X-ray diffraction to identify the crystal structure.
[0090]
The volume fraction of dispersed particles was determined as follows.
[0091]
That is, (1) The extraction residue of the “casting + hot working + heat treatment” material is quantified to determine the total mass ratio of dispersed particles present in the steel. (2) Next, the extraction residue is subjected to X-ray diffraction, the diffraction peak of each particle is specified, and the integrated intensity of each is obtained. (3) The mass ratio of each dispersed particle is determined from the integrated intensity ratio. (4) The specific gravity of the particles is determined from the composition of each dispersed particle determined in advance by EDX analysis. (5) The volume fraction is obtained using the specific gravity of each particle.
[0092]
Tables 3 and 4 also show the volume fractions of the hexagonal boride and the cubic carbonitride obtained as described above. In Tables 3 and 4, hexagonal boride is XB. ₂ , Cubic carbonitride is described as X (C, N), and XB ₂ The volume fraction of dispersed particles other than X (C, N) is also shown.
[0093]
[Table 3]

[0094]
[Table 4]

[0095]
The Young's modulus was obtained by measuring the resonance frequency by a normal transverse resonance method using a Young's modulus measurement sample having a thickness of 1.5 mm, a width of 10 mm, and a length of 60 mm.
[0096]
Tables 3 and 4 show the average Young's modulus obtained for the outer periphery, the center, and the intermediate portion between the outer periphery and the center for each of the “casting + heat treatment” material and the “casting + hot working + heat treatment” material. The values are organized and shown.
[0097]
Using a No. 14A test piece (diameter of 8 mm and rating distance of 56 mm) described in JIS Z 2201, a tensile test was performed at normal temperature and a strain rate of 1 / second. In Table 3 and Table 4, the average value of 0.2% yield strength calculated | required about the outer peripheral part, the center part, and the intermediate part of an outer peripheral part and a center part is arranged and shown.
[0098]
From Tables 3 and 4, the following is clear.
[0099]
Among steels 1-4 that do not contain C, TiB ₂ In the case of steel 1 and steel 2 whose volume fraction is less than 10%, primary TiB ₂ The maximum particle size of 7.0 and 16.4 μm is 20 μm or less, respectively, but the Young's modulus is low and the target 240 GPa is not reached. On the other hand, in the case of steel 3 and steel 4, TiB ₂ Although the volume fraction of is over 10%, the primary TiB ₂ Since the maximum particle size of the particles exceeds 20 μm, the Young's modulus is still low and does not reach the target of 240 GPa. Furthermore, the Young's modulus of the “casting + hot working + heat treatment” material is 4-7 GPa lower than the Young's modulus of the “forging + heat treatment” material. In addition, the 0.2% proof stress of the said steels 1-4 is 299 MPa at most, and its intensity | strength is low.
[0100]
Among steels 5 to 8 that do not contain B, in the case of steel 5 and steel 6 in which the volume fraction of Ti (C, N) is less than 10%, the largest grain of primary crystal Ti (C, N) The diameters are 4.0 μm, 19.0 μm, and 20 μm or less, respectively, but the Young's modulus is low and the target 240 GPa is not reached. In the case of steel 7 and steel 8, the volume fraction of Ti (C, N) is 10% or more, but the maximum particle size of primary crystal Ti (C, N) exceeds 20 μm. After all, Young's modulus is low and the target 240 GPa is not reached. In particular, the Young's modulus of the material subjected to hot working, that is, the Young's modulus of the “casting + hot working + heat treatment” material is 3-7 GPa lower than the Young's modulus of the “forging + heat treatment” material. In addition, the 0.2% proof stress of the said steel 5-8 is 552 MPa at most.
[0101]
In the case of steel 9, primary TiB ₂ And the maximum particle size of Ti (C, N) is 20 μm or less, and TiB ₂ And Ti (C, N) are included in a composite, but the total volume fraction is less than 10%, so the Young's modulus is low and the target 240 GPa is not reached.
[0102]
On the other hand, in the case of steel 10-20, the volume ratio XB specified in the present invention ₂ And X (C, N) are secured, and the primary crystal XB ₂ And the maximum particle size of X (C, N) is 20 μm or less, and a Young's modulus exceeding 240 GPa is obtained.
[0103]
Specifically, in the case of steel 10-12, XB ₂ And X (C, N) in a total volume fraction of 13 to 21.7%, and the primary crystal XB ₂ And the maximum particle size of X (C, N) is 11.0 to 14.8 μm and 20 μm or less, and a high Young's modulus of 248 to 275 GPa is obtained. Even when subjected to hot working, Young's modulus is hardly lowered.
[0104]
In the case of steel 13, since Nb is included, the 0.2% proof stress is increased to 720 MPa.
[0105]
In the case of steel 14, by containing a large amount of V, (Ti, V) B ₂ In addition, fine (Ti, V) (C, N) are dispersed at a high density, and an extremely high strength of 998 MPa is obtained with a 0.2% proof stress.
[0106]
In the case of steel 15, by containing a large amount of Nb, (Ti, Nb) B ₂ In addition, fine (Ti, Nb) (C, N) is dispersed at a high density, and a high strength of 886 MPa is obtained with a 0.2% proof stress.
[0107]
In the case of steel 16, since V and Nb are contained in combination, a high strength of 970 MPa is obtained with a 0.2% proof stress.
[0108]
In the case of steel 17, the impurity element is fixed by containing Nd, and the adhesion between the dispersed particles and the Fe matrix is improved, so that the Young's modulus after hot working is improved.
[0109]
In the case of steel 18, by containing a large amount of V, an extremely high strength of 1011 MPa is obtained with a high Young's modulus and 0.2% proof stress.
[0110]
In the case of steel 19, since Ca and Nd are combined and contained, Young's modulus is improved after hot working.
[0111]
Steel 20 is a hard steel containing a large amount of V, but contains Mg. Therefore, even if the finished diameter of the bar material by hot working is made smaller, that is, even when the total area reduction is large, it is Young. The rate has not decreased.
[0112]
In the case of steels 21 to 23, the amount of particles dispersed in the matrix is large, but a high Young's modulus cannot be obtained.
[0113]
That is, in the case of steel 21, TiB ₂ And Ti (C, N) have a volume fraction of 25.9%, but primary TiB ₂ And the maximum particle size of Ti (C, N) exceeds 20 μm and is 24.2 μm, the Young's modulus is low and the target of 240 GPa is not achieved.
[0114]
In the case of steel 22, TiB ₂ And Ti (C, N) total volume fraction is as low as 5.2%, and 15% of the particles dispersed in the Fe matrix are Fe ₃ C (cementite), Fe ₃ Since the Young's modulus of C is lower than that of the Fe matrix, the Young's modulus is extremely low.
[0115]
In the case of steel 23, TiB ₂ And Ti (C, N) total volume fraction is as low as less than 2.0% and does not contribute to high Young's modulus ₃ Since it contains 9% C and 15% orthorhombic boride, the Young's modulus is extremely low.
[0116]
【The invention's effect】
The high-strength, high-rigidity steel of the present invention can suppress the coarsening of crystallized particles even if it is manufactured by casting molten steel, can be formed by hot working, and has a Young's modulus of 240 GPa or more. Therefore, for example, it can be used for various structural parts for automobile engines, large structural members such as bolts and various shafts.

Claims (4)

A high-strength, high-rigidity steel produced by casting molten steel, in which a hexagonal boride and a cubic carbonitride are dispersed in a total volume of 10 to 35% in an Fe-based alloy matrix, and A high-strength, high-rigidity steel, wherein the hexagonal boride and the cubic carbonitride have a particle size of 20 μm or less. The hexagonal boride is one or more of TiB ₂ , VB ₂ , NbB ₂ and composite borides thereof, and the cubic carbonitride is Ti (C, N), V (C, N ), Nb (C, N), Ta (C, N), Hf (C, N) and one or more of these composite carbonitrides, the high strength according to claim 1 High rigidity steel. In mass%, C: 0.2-1.7%, Si: 0.05-0.5%, Mn: 0.2-1.5%, Ti: 1-12%, Mo: 0.1 2%, B: 0.5-3.2%, N: 0.001-0.01%, V: 0-12%, Nb: 0-12%, Ta: 0-1%, Hf: 0 1%, Ca: 0 to 0.007%, Mg: 0 to 0.007% and Nd: 0 to 0.007%, with the balance being Fe and impurities, represented by the following formula (1) The high-strength, high-rigidity steel according to claim 2, wherein SB exceeds 100 and less than 1050, and SC represented by the following formula (2) exceeds 50 and less than 550.

However, the element symbol in each formula represents the content in mass% of the element. In mass%, C: 0.2-1.7%, Si: 0.05-0.5%, Mn: 0.2-1.5%, Ti: 1-12%, Mo: 0.1 2%, B: 0.5-3.2%, N: 0.001-0.01%, V: 0-12%, Nb: 0-12%, Ta: 0-1%, Hf: 0 1%, Ca: 0 to 0.007%, Mg: 0 to 0.007% and Nd: 0 to 0.007%, with the balance being Fe and impurities, represented by the above formula (1) In the casting process of the molten steel in which the compound of each element in the steel in which the SB exceeds 100 and less than 1050 and the SC represented by the formula (2) exceeds 50 and less than 550 is completely dissolved in the steel, Hexagonal borides or cubic carbonitrides are crystallized, and then the three phases of hexagonal boride, cubic carbonitride and Fe matrix phase are crystallized from the residual molten steel. The steel ingot is cooled to a temperature range of 300 ° C. or less, and further homogenized at a temperature of HT (° C.) or less represented by the following formula (4). Perform hot working so that the total area reduction is 50% or more at temperature, and then perform normalization at 850 to 1050 ° C., cool to 300 ° C. or lower, and then temper at 700 ° C. or lower. A method for producing high-strength, high-rigidity steel.

However, the element symbol in the formula (4) represents the content in mass% of the element.