JP2004060049A - Ultrahigh temperature hot forged non-heat-treated part and method for producing the same - Google Patents
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Abstract
Description
【０００１】
【産業上の利用分野】
本発明は、自動車や建設機械の足廻りに使用される部品の中で、高強度・高靭性を必要とされる熱間鍛造品において、従来の熱間鍛造部品より複雑形状な部品を製造したり、製造時の材料歩留まり向上しつつ、鍛造後の焼入れ焼戻しを行うことなく高強度、高靭性を確保することを可能とした超高温熱間鍛造非調質部品とその製造方法に係るものである。
【０００２】
【従来の技術】
【特許文献１】特開平１−１９８４５０号公報
【特許文献２】特開平５−１５９３５号公報
従来、自動車部品、建設機械部品の中で、高強度，高靭性を必要とする熱間鍛造部品は、熱間鍛造後に調質、即ち焼入れ焼戻しして製造されていた。しかし、製造コストに占める調質コストが大きいことから非調質化が進められ、特許文献１で開示される様に、熱間鍛造後に放冷ままで強度，靭性を確保できる熱間鍛造用非調質鋼が開発された。ところが、さらなる製造コスト低減から、熱間鍛造時の材料歩留まり向上が要望されている。また、自動車の軽量化から部品の小型化が指向されている。その為には部品剛性を確保する必要があり、従来の熱間鍛造より複雑形状な部品となり、その結果、既存の鍛造機で成形できないような鍛造荷重増大を招くことになる。これらを解決するには、熱間鍛造時の鋼材の変形抵抗を低減することが必要である。これに対し、特許文献２では、１１５０〜１２５０℃とされる従来の熱間鍛造温度より高い温度で鋼材を加熱して熱間鍛造する超高温熱間鍛造方法が開示されている。従来の熱間鍛造が１１５０〜１２５０℃とされるのに対し、特許文献２記載の超高温熱間鍛造での加熱温度は、その下限を固相線温度の４５℃下とし、上限を液相線温度の２０℃下としている。しかし、そこまで鋼材を加熱するとオーステナイト粒が粗大化し、鍛造後放冷ままでは靭性が確保できない欠点を有する。その為、特許文献２では、超高温熱間鍛造直後に表層急冷を行うなどしている。
【０００３】
【発明が解決しようとする課題】
本発明は、高強度，高靭性を必要とされる熱間鍛造品において、従来の熱間鍛造部品より複雑形状な部品を製造したり、製造時の材料歩留まりを向上させるための超高温熱間鍛造において、焼入れ焼戻しすることなく鍛造後の靭性劣化などの問題を解決するものである。
【０００４】
【課題を解決するための手段】
（１）質量％で、
Ｃ：０．１〜０．６％、 Ｓｉ：０．２〜２．０％、
Ｍｎ：０．５〜２．５％、 Ｓ：０．０２〜 ０．１０％、
Ｃｒ：０．１〜１．０％、 Ｖ：０．０３〜０．３％、
Ａｌ：０．００２〜０．０６％、 Ｎ：０．００３〜０．０２％
を含有し、残部Ｆｅおよび不可避不純物からなり、組織はフェライトとパーライトからなり、フェライト分率０．１〜０．６、フェライト平均粒径１０〜３０μｍ、ＪＩＳ Ｇ ０５８８で 規定する全脱炭層深さＤＭ−Ｔが０．０２〜０．０８ｍｍであることを特徴とする超高温熱間鍛造非調質部品、
（２）質量％で、（１）に記載の超高温熱間鍛造非調質部品に加えて、更に
Ｔｉ：０．００３〜０．０５％、 Ｍｇ：０．０００２〜０．００５％、
Ｚｒ：０．０００２〜０．００５％
の中から１種または２種以上を含有することを特徴とする超高温熱間鍛造非調質部品、
（３）引張強さが８００〜１０００ＭＰａであることを特徴とする（１）又は（２）記載の超高温熱間鍛造非調質部品、
（４） （１）〜（３）の何れか１項に記載の非調質部品を製造する方法であって、（１）又は（２）記載の成分からなる鋼材を、下限温度を平衡状態図における固相線温度×０．９４または１２５０℃の何れか高い方とし、上限温度を固相線温度とする範囲に加熱し、前記範囲の温度域で超高温熱間鍛造加工した後、さらに加工品を９００℃以上、１２５０℃以下の温度で対数ひずみで０．２以上の仕上げ鍛造を加えた後放冷することを特徴とする超高温熱間鍛造非調質部品の製造方法、
（５） （１）〜（３）の何れか１項に記載の非調質部品を製造する方法であって、（１）又は（２）記載の成分からなる鋼材を、固相線温度以上、加熱上限温度Ｔｕ［℃］以下の加熱限界域に加熱し、前記加熱限界域における保持時間を１分以下とし、そのまま下限温度を固相線温度×０．９４または１２５０℃の何れか高い方とし、上限温度を加熱上限温度Ｔｕ［℃］以下とする範囲で超高温熱間鍛造加工した後、さらに加工品を９００℃以上、仕上げ上限温度Ｔｆ［℃］以下の温度で対数ひずみで０．２以上の仕上げ鍛造を加えた後放冷することを特徴とする超高温熱間鍛造非調質部品の製造方法。
ただし、Ｔｕ＝１５３６−（３１７．０×Ｃ％＋９．４×Ｓｉ％＋５．２×Ｍｎ％＋１４０．３×Ｓ％）
Ｔｆ＝１４５６−（４４０．０×Ｃ％＋１２．６×Ｓｉ％＋１２．８×Ｍｎ％＋３５．３×Ｓ％）
であり、Ｃ％、Ｓｉ％、Ｍｎ％、Ｓ％は、質量％で表したＣ、Ｓｉ、Ｍｎ、Ｓの含有量である。
（６） 超高温熱間鍛造において素材表面の９０％以上が金型に接触する様に超高温熱間鍛造加工することを特徴とする（４）又は（５）記載の超高温熱間鍛造非調質部品の製造方法。
（７） 超高温熱間鍛造加工を行った後、鍛造機の下死点において少なくとも材料表面温度が９００℃以上、１２００℃以下となるまで鍛造時の最大荷重の１０〜８０％の荷重で保持することを特徴とする請求項（４）〜（６）の何れか１項に記載の超高温熱間鍛造非調質部品の製造方法。
【０００５】
【発明の実施の形態】
本発明の根幹をなす技術思想は以下の通りである。
【０００６】
本発明では鍛造放冷ままで所定の強度、靭性が得られるように成分調整された熱間鍛造非調質鋼を用いるが、非調質鋼を用いて特許文献２記載の様に単に超高温鍛造したままでは、必要な靭性が得られない。しかし、熱間鍛造は通常一回の加工だけでなく、何工程かの鍛造加工を受けて所定の形状に成形される。そこで、本発明では、粗成形を超高温熱間鍛造で行い、続けて行われる仕上げ成形を９００℃〜１２５０℃の温度で行うこととした。即ち、従来の熱間鍛造より複雑形状に成形するため、超高温熱間鍛造の粗成形でほぼ製品形状に成形する。そして、そのままでは高靭性を確保できないことから、所定の温度、歪みの仕上げ成形を行うことで靭性を確保するのである。また、鍛造後の表面フェライト脱炭層が所定の範囲にあることで、固相線近くまで加熱しても表層フェライトが脱炭しているので、表面、特に角部での溶融割れを防止することが可能となる。
【０００７】
以下に本発明を詳細に説明する。
【０００８】
Ｃは、鋼を強化するのに有効な元素であるが、０．１％未満では充分な強度が得られない。一方、過多に添加すると靭性が低下するため、添加量の上限を０．６％とする。
【０００９】
Ｓｉは、脱酸材として働き、固溶強化元素として使われる。０．２％未満では脱酸材としての作用が不足し、過多に添加すると必要以上に強度を上げて靭性を低下させるため、添加量の上限を２．０％とする。
【００１０】
Ｍｎは、強度の調整と脱酸作用をする。０．５％未満では強度が不足し、２．５％を越えると靭性が低下するとともに、熱間圧延時に割れが生じて製造が困難となる。
【００１１】
Ｓは、被削性を向上させるのに不可欠な元素であり、その効果は０．０２％より多いＳで期待される。しかし、０．１０％を越えると靭性を低下させる。
【００１２】
Ｃｒは、Ｍｎと同様に強度を補うための元素であり、その効果を得るには０．１％以上添加する必要がある。強度を補うために１．０％まで添加させることができるが、１．０％を越えると靭性を劣化させる。
【００１３】
Ｖは、固溶強化、析出強化により靭性を向上させる。この効果を得るには０．０３％以上の添加が必要である。しかし、過多に添加しても効果の向上が小さく、むしろ靭性を低下させるのでその上限を０．３０％とした。
【００１４】
Ａｌは、鋼の脱酸および結晶粒の微細化のために有効な元素であるが、０．００２％未満ではその効果がない。一方、過多に添加すると靭性を低下させるために添加量の上限を０．０６％とする。
【００１５】
Ｎは、Ｖ炭窒化物を生成し析出強化のために必要な元素であるが、０．００３％未満では充分な効果が得られない。一方、過多に添加すると固溶したＮによって靭性が劣化するため、添加量の上限を０．０２％とする。
【００１６】
Ｔｉは，窒化物・炭化物を生成する。窒化物は高温まで固溶せずに残るため、加熱時のオーステナイト粗大化を防止するのに有効である。また炭化物は微細に分散して析出強化に有効である。０．００３％未満ではこれらの効果は現れず、０．０５％を越えると靱性が劣化するため、その添加量の下限を０．００３％、上限を０．０５％とする。
【００１７】
ＭｇおよびＺｒはともに酸化物や硫化物、あるいはこれらの複合物を形成し、加熱時のオーステナイトの粗大化を抑制する効果を持つ元素である。また、これらの酸化物はＭｎＳの析出核になるため被削性も向上する。いずれも、０．０００２％未満ではその効果はなく、０．００５％を越えると、靱性が劣化するため、添加量の上限を０．００５％とする。
【００１８】
鍛造後の冷却速度によって組織はフェライト＋パーライト、ベイナイト、マルテンサイトに区別できるが鍛造後放冷ままで高強度、高靭性を得られる組織として本発明ではフェライト＋パーライト組織とした。
【００１９】
フェライト分率が多いほど靭性は向上するが、０．６を越えたフェライト分率は強度低下を招く。また、フェライト分率が低いと靭性が低いことからその分率を０．１〜０．６とした。そしてフェライト分率は、鋼材のＣ量，Ｓｉ量を適正に添加することで制御が可能となる。なお、フェライト分率は、光学顕微鏡により２００〜５００倍で３〜５視野観察し、フェライトとパーライトのコントラストを二値化して解析システムでその分率を算定した値と定義する。
【００２０】
フェライト平均粒径は、仕上げ温度を低くし、強加工をした方が小さくできるが、これは結果として金型負荷の急増を招き、機械的性質の向上のために必要以上に粒径の目標値を小さくすると、鍛造自体が困難となる。従って、フェライト平均粒径は実際に製造可能なレベルを考えて、その下限を１０μｍに設定した。またフェライト平均粒径が３０μｍを越えると靭性劣化を招くため、フェライトの平均粒径を１０〜３０μｍとした。フェライト平均粒径の制御には、鋼材に添加するＡｌ量の最適化とともに，Ｔｉ，Ｍｇ，Ｚｒ等の添加が重要である。また、鋼材成分のみならず超高温鍛造後に行う仕上げ鍛造が必要であり鍛造温度，加工量の制御が重要である。特にフェライト平均粒径を小さくするには仕上げ鍛造条件として、後述する様に９００〜１２５０℃の温度域で対数ひずみで０．２以上の加工を行うことが必要となる。なお、フェライト平均粒径は光学顕微鏡により２００〜５００倍で３〜５視野観察し、切断法により求めた値と定義する。
【００２１】
さらに、鍛造後の全脱炭層深さＤＭ−Ｔであるが、０．０８ｍｍを越える脱炭層は鍛造品強度を著しく低下させる。特にシャフト部を有する熱間鍛造品では、シャフト部に捻り変形を受けることが多く表面強度が重要となる。脱炭は超高温熱間鍛造の加熱段階で発生するが、全脱炭層深さが０．０２ｍｍより少なくなるような高速加熱では、加熱時に素材表面、特に熱集中を受ける角部等で溶融割れが発生するという問題がある。全脱炭層深さを０．０８ｍｍ以下にするには、高周波などによる急速加熱が必要となる加熱装置にもよるが、周波数等を制御する、或いは加熱装置の中を移動する素材の送り速度を制御する等の素材加熱時間を制御することで全脱炭層深さを制御することが可能となる。
【００２２】
引張強さは、鍛造品の軽量化の点で下限を８００ＭＰａに限定した。一方、１０００ＭＰａを越えると、靭性が著しく低下し、切削寿命および金型寿命も著しく低下するため、上限を１０００ＭＰａ以下にした。
【００２３】
次に、製造方法について述べる。
【００２４】
加熱温度は超高温熱間鍛造によって変形抵抗を低減させるために、固相線温度×０．９４または１２５０℃の何れか高い方を下限とした。これより低い温度では充分に変形抵抗が低くならず、材料流動が十分に行われないからである。上限温度を固相線温度とするのは、固相線を越える温度では結晶粒界などが溶融するからである。素材内部が溶融すると、内部鍛造時にこの溶融部分が空孔化しやすくなるため、空孔を低減できる条件での鍛造が必須となる。従って本発明では、上限温度を固相線温度とした。
【００２５】
固相線温度は、鉄と鋼７３巻４号Ｓ１９６頁１９８７年に記載される析出物の凝固過程を観察するのに用いられる一方向凝固実験によって推定される。高周波加熱とカーボンサセプターを用いて炉内に温度勾配を持たせ、その炉内で棒材を加熱し、その後急冷する。急冷した棒材の内部組織観察から棒材各位置での温度と組織を対応させて素材の固相線温度を推定した。なお、一方向凝固実験（ｕｎｉｄｉｒｅｃｔｉｏｎａｌ Ｓｏｌｉｄｉｆｉｃａｔｉｏｎ ｅｘｐｅｒｉｍｅｎｔ）の詳細は、松宮（Ｍａｔｓｕｍｉｙａ）、他４名「連続鋳造スラブにおける偏析の数値解析（Ｍａｔｈｅｍａｔｉｃａｌ Ａｎａｌｙｓｉｓ ｏｆ Ｓｅｇｒｅｇａｔｉｏｎｓ ｉｎ Ｃｏｎｔｉｎｕｏｕｓｌｙ−ｃａｓｔ Ｓｌａｂｓ．） 」、日本鉄鋼協会論文集（Ｔｒａｎｓａｃｔｉｏｎｓｏｆ Ｔｈｅ Ｉｒｏｎ ａｎｄ Ｓｔｅｅｌ Ｉｎｓｕｔｉｔｕｔｅ ｏｆ Ｊａｐａｎ）、日本鉄鋼協会、１９８４年１１月、２４巻、１１号、ｐ．８７５、８７６に記載されている。
【００２６】
超高温熱間鍛造は、加熱と同じ温度域で行う。固相線温度×０．９４または１２５０℃の何れか高い方より低い温度では変形抵抗が高く、金型寿命が低下するためと、材料流動が悪いため、複雑形状部品へ成形できない。また、固相線温度より高い温度では、鍛造時の加工発熱で素材内部に溶融部が発生し、靭性劣化を招くからである。
超高温熱間鍛造で粗成形した後に仕上げ鍛造を行う。この仕上げ鍛造は、製品形状を所定の形状にすると共に、本発明では靭性付与を目的としており、対数ひずみで０．２以上の鍛造加工を行わないと靭性付与の効果が小さいからである。なお、超高温熱間鍛造でほぼ製品形状に成形することから仕上げ鍛造での加工率は対数ひずみで０．５〜２．０位が望ましい。なお、実際の歪は部位により一様ではなく不均一な歪分布となるが、優れた靱性を確保できる範囲を明確にするためには、決められた式で加工度の大小を表現する必要がある。そこで本発明は鍛造時の対数ひずみは加工前後素材形状から算出し、据込み部分であれば、対数ひずみはｌｎ（加工前の素材高さ／加工後の素材高さ）で求め、押出し部分の対数ひずみはｌｎ（加工前の素材断面積／加工後の素材断面積）で求めるというようにした。この式により求めた値が０．２以上であれば優れた靱性を確保することができる。
仕上げ鍛造を１２５０℃以下で行うのは、これより高い温度では鍛造後のオーステナイト粒が大きく靭性が低いからである。靭性を向上させるには低温で加工することによってフェライト分率を上げることが望ましいが、９００℃より低い温度では変形抵抗が急増するため金型寿命の低下を招くからである。
なお固相線温度以上に鋼材を加熱しても、加熱上限温度をＴｕ［℃］以下とし、加熱時間が短ければ、結晶粒界などの溶融の防止が可能である。加熱上限温度Ｔｕ［℃］は、
Ｔｕ＝１５３６−（３１７．０×Ｃ％＋９．４×Ｓｉ％＋５．２×Ｍｎ％＋１４０．３×Ｓ％）
により求めることができる。ここで、、Ｃ％、Ｓｉ％、Ｍｎ％、Ｓ％は質量％で表したＣ、Ｓｉ、Ｍｎ、Ｓの含有量である。
鋼材が固相線温度以上、加熱上限温度Ｔｕ［℃］以下の加熱限界域に曝される保持時間は、１分以内であれば、内部に空孔を生じることなく超高温熱間鍛造することが可能であり、オーステナイト粒の粗大化を抑制することができる。鋼材を固相線温度に加熱し、直ちに超高温熱間鍛造を行っても良いため、保持時間の下限は０分で良い。
このような加熱を行う際には、加熱速度を大きくすることが好ましく、高周波加熱又は通電加熱が好適である。また、加熱、保持する際には、高周波コイルと鋼材の間に送風を行う等、鋼材の表層のみを冷却し、表層部の溶融を防止することが好ましい。
また、超高温鍛造加工前のオーステナイト粒の粗大化を防止したことにより、超高温鍛造加工後のオーステナイト粒径の粗大化も抑制される。したがって、仕上げ鍛造温度の上限を、１２５０℃超としても、仕上げ上限温度Ｔｆ［℃］以下であればフェライト平均粒径が微細になり、鍛造後の靱性を確保することが出来る。仕上げ上限温度Ｔｆ［℃］は、
Ｔｆ＝１４５６−（４４０．０×Ｃ％＋１２．６×Ｓｉ％＋１２．８×Ｍｎ％＋３５．３×Ｓ％）
によって求めることができる。ここで、Ｃ％、Ｓｉ％、Ｍｎ％、Ｓ％は、質量％で表したＣ、Ｓｉ、Ｍｎ、Ｓの含有量である。
【００２７】
本発明の目的の一つとして、材料歩留まりの向上がある。その為に、超高温熱間鍛造での材料流動向上によっても材料歩留まりは向上できるが、その際密閉金型、或いは半密閉金型を用いることで製品形状に近い形まで成形でき、材料歩留まりは格段に向上される。従来の熱間鍛造では変形抵抗が高く密閉度が高い金型を用いても金型内へ材料が充満しない。しかし超高温熱間鍛造では変形抵抗が低いことから、密閉度が高い金型を用いても金型内への材料充満が良くなる。しかし、素材表面の９０％以上が金型に接触する様に超高温熱間鍛造しないと、材料歩留まり向上効果が小さく、製品形状がでにくくなる。なお、高い密閉度の金型で超高温熱間鍛造成形する方が材料歩留まりは向上するが、同時に鍛造荷重が急増するので、９８％以下が望ましい。さらに、密閉度の高い金型を使うのでなく、既存金型を用いて鍛造機のラムの下死点を下げて使用する素材のサイズを小さくすることでも、接触率はあがり素材歩留まりは向上できる。
【００２８】
仕上げ鍛造を所定の温度で行うためには、超高温熱間鍛造加工した後に素材温度が下がるまでの時間を必要とする場合がある。この場合、仕上げ鍛造するまでの間に熱ひずみにより製品形状の精度が悪化する。そこで超高温熱間鍛造加工を行った後に形状を凍結しておくことで製品精度が向上される。鍛造機の下死点で素材の表面温度が９００〜１２００℃となるまで保持する。表面温度が１２００℃以下となるまで荷重保持するのは熱ひずみを小さくするためである。なお、好ましくは１０００℃以下となるまで荷重保持することが望ましい。下限の温度を９００℃以上とするのは、９００℃より低くなると次の仕上げ鍛造で鍛造荷重が急増し、金型寿命が低下するからである。また、保持荷重を鍛造時の最大荷重の１０％以上とするのは、これ未満では熱ひずみの抑制が不十分であるからである。なお、保持荷重が大きくなると金型寿命へ影響するため、その上限は最大荷重の８０％以下とする。なお、鍛造機の下死点で荷重保持することに代えて、水などで表層冷却することも可能である。
【００２９】
【実施例】
表１のＡ〜ＡＡに示す化学成分を有する鋼材を用いて本発明例と比較例の実験を行った。鋼種Ａ〜Ｑが本発明例の対象鋼種であり、鋼種Ｒ〜ＡＡが比較例に用いられた鋼種である。なお同表には、参考として隣（Ｐ）の成分と、固相線温度（Ｔｓ）も併記する。固相線温度はφ１５×２５０ｍｍの棒状素材を一方向凝固試験から推定した温度である。
【００３０】
表２は、表１に示した鋼種を用いて超高温熱間鍛造−仕上げ鍛造したの後の強度，靭性を調べた結果を示す。実験はφ６０×６０ｍｍの素材を所定の温度に高周波で加熱した。なお加熱時の周波数は３〜５ＫＨｚであり、室温から１２５０℃までを５℃／秒の速度で加熱しその後は１℃／秒で所定の温度まで加熱した。所定の温度になってからは約３０秒間保持した後鍛造に供した。高周波加熱後に鍛造機まで試料を移動するため素材温度が低下する。そこで加熱温度をＴｋ（℃）、鍛造直前温度をＴｔ（℃）として表２に示した。鍛造直前温度Ｔｔは、素材表面を放射温度計で測温した結果である。鍛造は油圧サーボ機構を有する圧縮試験機にて、ラム速度２００ｍｍ／ｓで圧縮試験を行った。超高温熱間鍛造はφ６０×６０ｍｍの試料を横置きし、平坦な圧盤を用いて圧縮率約５０％の圧縮を行った。加工前の高さ６０ｍｍの素材を３０ｍｍへ加工することで圧縮率は５０％となり、対数ひずみはｌｎ（加工前の素材高さ／加工後の素材高さ）で算出され、この場合の対数ひずみは０．７（＝ｌｎ（６０／３０））となる。また、超高温鍛造後にさらに仕上げ鍛造を行うが、超高温鍛造で高さ３０ｍｍとした素材をさらに１６ｍｍへ加工する場合、仕上げ鍛造での対数ひずみは０．６（＝ｌｎ（３０／１６））となる。
【００３１】
超高温鍛造後仕上げ鍛造するまでには試験機の油圧が回復するまでに時間を必要とし、超高温熱間鍛造時の温度より低くなるため、仕上げ直前の素材表面温度を放射温度計にて測温した結果を表２に示している。仕上げ鍛造後は放冷した後、フェライト分率、フェライト平均粒径、全脱炭層深さを測定した。脱炭層深さは光学顕微鏡での倍率を２００〜５００倍として、読み取り寸法を付した接眼鏡を用いて脱炭層深さを１０カ所測定し、最大、最小を除く８カ所の平均値とした。また、圧縮後の試料から引張試験片、シャルピー衝撃試験片を採取した。表２に引張強さ、衝撃値を示すが衝撃値は−５０℃での吸収エネルギーである。
【００３２】
表２から、本発明例１〜１７は引張強さが約８００〜９００ＭＰａであり、衝撃値も２０Ｊ／ｃｍ２以上となっている。一方、比較例１〜１０は表１の比較例Ｒ〜ＡＡの鋼種を用いた場合であり、いずれも衝撃値が２０Ｊ／ｃｍ２にいたっていない。特に比較例５では強度も７８７ＭＰａと低い。比較例１１は仕上げ鍛造前温度が８５０℃と低いため、フェライト分率も所定より高くなり強度が低い。比較例１２ではさらに仕上げ前温度が低い場合であり、この様な低い温度で仕上げ鍛造するとフェライト粒径が所定より小さくなり、本発明に比べ高強度を得ることができるが、仕上げ鍛造時の荷重が約３０００ＫＮと高く、本発明の約７００ＫＮに対して４倍以上の荷重となっており、金型、鍛造機への負担が急増することになる。比較例１４は仕上げ鍛造時の加工が小さい場合であるが、いずれも鍛造後の靭性が低かった。比較例１５は超高温熱間鍛造時の加熱温度、及び鍛造直前温度が高い場合であり、脱炭層が深くなり所定の強度に達していなかった。また、比較例１５では、加熱時に試料の一部の溶融も見られた。比較例１６は加熱後の保持温度時間を５分間と長くした場合であり、この場合も脱炭増が深く強度が低かった。
【００３３】
表３は、さらに加工条件を種々代えた場合の比較である。本発明例では超高温熱間鍛造時の荷重は２８２〜４１６ＫＮ、仕上げ鍛造時では７０１〜７９６ＫＮである。一方、比較例１７は加熱温度が低く、比較例１８は所定の温度に加熱するものの超高温熱間鍛造までの時間を長くすることで超高温熱間鍛造直前温度を所定より低くした場合であるが、いずれも６００ＫＮ以上と高い荷重となっている。比較例１９は仕上げ鍛造時直前温度が低い場合であるが、このときの仕上げ鍛造荷重は１８２６ＫＮと本発明例での仕上げ鍛造時に比べて倍以上となっている。鍛造荷重が高いことは金型寿命の低下を招くことになる。比較例２０は仕上げ鍛造を行わなかった場合であるが靭性値が低いことが分かる。比較例２１は脱炭層深さを浅くするため、所定の温度に加熱した後に３０秒の保持を行わずに直ぐに超高温熱間鍛造を行った場合である。温度測定は試料の表面で行っており、表面温度は所定の温度に達しているものの、温度保持を行わないため中心部の素材は所定の温度にまで加熱されていないためか、熱間鍛造時の荷重が８６４ＫＮと本発明での超高温熱間鍛造時の荷重より高くなってしまった。
【００３４】
表４は、前記（７）に係る発明の効果を確認するため、超高温熱間鍛造後の試料に圧縮機のラムで所定の荷重を保持した場合である。いずれの場合も所定の条件で１０個の試料を鍛造して仕上げ鍛造後の鍛造品の高さを測定した。１０個中の最大値の高さの物と最小高さの物の高さの差を同表のＨｂで示している。なお、加熱温度，鍛造前温度，荷重などは１０個中の一つの試料において測定した結果である。本発明例ではＨｂは０．２４〜０．３４ｍｍであった。一方、参考例１は保持荷重が小さく、参考例２は保持温度が高い場合であり、Ｈｂは０．６２ｍｍ以上と前記（７）の条件を満たす本発明例よりかなり鍛造後試料の高さのバラツキが多かった。なお、参考例３は保持温度が低い場合であり、Ｈｂは小さいものの保持温度を低くすると仕上げ鍛造時の鍛造前温度も低くなり仕上げ鍛造時の荷重が倍近く高くなってしまい金型寿命を悪化させてしまう。
【００３５】
表５は前記（６）に係る発明の効果を確認するために、鍛造歩留まりを検討すべく図１の様な部品形状にて鍛造を行った。鍛造機は一般のクランクプレスを用い、加熱は高周波炉にて行った。同表に示すサイズの丸棒を用い、超高温熱間鍛造後、仕上げ鍛造を行い、図１の形状にトリミングを行った。トリミング後の鍛造品の重量を量ることにより、鍛造前前素材との重量比から素材歩留まりを求めた。超高温熱間鍛造での素材−金型の接触率は、鍛造機ラムの下死点と金型を上下させることで調整した。鍛造機のまた鍛造品のアーム部から引張試験片，衝撃試験片を採取した。アーム部の加工量であるが、アーム部は超高温熱間鍛造において直径５３〜５８ｍｍの素材を高さ３３ｍｍまで加工した。その後、仕上げ鍛造で１８ｍｍにしており、仕上げ鍛造での対数ひずみは０．６（＝ｌｎ（３３／１８））となる。衝撃値は−５０℃での値で代表した。まず実験として通常の熱間鍛造に相当する比較例２２の実験を行った。φ５８×１７７ｍｍの素材で接触率９２％で行うことで未成形部が発生しないように加工でき、その時の素材歩留まりは６２％であった。次に超高温熱間鍛造での実験を行ったものが本発明例２４，本発明例２５である。本発明例２４は比較例２２よりも小さいサイズの試料で接触率９３％でも未成形部無く加工できた。これは比較例２２の熱間鍛造は加熱温度が低いため材料流動が悪く、金型フラッシュ部から十分にバリを出さないと接触率９２％を確保できないからである。一方本発明例２４では、材料流動が良いことからバリを出さなくても接触率を確保できた。その結果バリの量が少ない本発明例２４での歩留まりが格段に向上していた。本発明例２５は、鍛造機ラムの下死点を調整して接触率を向上させた場合であるが、歩留まりも８７％と向上している。参考例４，参考例５は下死点を調整して接触率を８５％にした場合であるが、本発明例２４と同じサイズの試料で鍛造した参考例４では、未加工部が発生した。そこで未加工部が発生しないようにするにはφ５５×１７４ｍｍの素材を用いることが必要である、それが参考例５であるが、素材歩留まりは７１％と低かった。以上から本発明により高い歩留まりで鍛造できることが分かり、安い製造コストで製造できることになる。
表６は、前記（５）に係る発明の効果を確認するため、鋼材を固相線温度以上、Ｔｕ［℃］以下に１分以内加熱し、仕上げ鍛造の上限温度をＴｆ［℃］以下とした場合の比較である。表１のＡ〜Ｄ鋼の成分からなるφ６０×６０ｍｍの試料を用い、表６に示した条件で、約５０％の超高温熱間鍛造を行った後、仕上鍛造し、放冷した。また、中心部に熱電対を取り付けた試料を作製し、表６に示した条件での超高温熱間鍛造及び仕上鍛造を別途実施し、超高温熱間鍛造前及び仕上鍛造前の中心温度を測定した。その結果を表６に併記する。
表６において、本発明例２６〜３３は、何れも溶融割れを生じることなく鍛造でき、鍛造後の−５０℃での衝撃値も２０Ｊ／ｃｍ２以上となっている。一方、比較例２３，２４は加熱温度が高く、加熱中に溶融し、超高温熱間鍛造が出来なかった。そのため、表中の超高温熱潤鍛造の鍛造前温度、鍛造荷重、仕上げ鍛造、鍛造後の製品の欄には”−”を記入した。比較例２５，２６は固相線温度以上に加熱される時間が長く、超高温熱間鍛造荷重は低いものの鍛造割れが生じた。そのため、仕上げ鍛造を行うことが出来ず、表中の仕上げ鍛造、鍛造後の製品の欄には”−”を記入した。比較例２７，２８は仕上鍛造温度が高いため、また比較例２９，３０は仕上鍛造のひずみが小さいため、鍛造後の−５０℃での衝撃値が２０Ｊ／ｃｍ２に達していない。
参考例は前記（４）の条件で実験した結果である。参考例６〜９での超高温熱間鍛造荷重が３８２〜４７７ｋＮ、仕上鍛造荷重が１８２６〜２２２５ｋＮであるのに対し、同等のひずみ条件である本発明例２６，２７，３２，３３では超高温熱間鍛造荷重が２１１〜２８５ｋＮ、仕上鍛造荷重が１４９５〜１５３９ｋＮと、大幅な荷重低減効果が認められる。
【００３６】
【表１】
【表２】
【表３】
【表４】
【表５】
【表６】
【発明の効果】
本発明により鋼材の成形性が著しく高まることにより、従来成し得なかった複雑形状部品の加工や、高い素材歩留まりで加工できる。このことは、部品の軽量化を実現するとともに、従来より高い生産性，安いコストでの製造を実現できることになり、機械部品の製造において多大の効果をもたらすものである。
【図面の簡単な説明】
【図１】鍛造実験における鍛造品形状と概略寸法を示す図。[0001]
[Industrial application fields]
The present invention manufactures parts that are more complex than conventional hot forged parts in hot forged parts that require high strength and high toughness, among parts used in the footings of automobiles and construction machinery. Or high temperature hot forged non-refined parts that can secure high strength and high toughness without quenching and tempering after forging while improving the material yield at the time of manufacturing, and its manufacturing method. is there.
[0002]
[Prior art]
[Patent Document 1] Japanese Patent Laid-Open No. 1-198450
[Patent Document 2] JP-A-5-15935
Conventionally, among automotive parts and construction machine parts, hot forged parts that require high strength and high toughness have been manufactured by tempering, that is, quenching and tempering after hot forging. However, since the tempering cost occupying the manufacturing cost is large, non-tempering has been promoted, and as disclosed in Patent Document 1, non-hot tempering for hot forging that can ensure strength and toughness while being allowed to cool after hot forging. Tempered steel was developed. However, in order to further reduce manufacturing costs, there is a demand for improved material yield during hot forging. In addition, miniaturization of parts is aimed at from the weight reduction of automobiles. For this purpose, it is necessary to ensure the rigidity of the parts, resulting in a part having a more complicated shape than the conventional hot forging, resulting in an increase in forging load that cannot be formed by an existing forging machine. In order to solve these problems, it is necessary to reduce the deformation resistance of the steel material during hot forging. On the other hand, Patent Document 2 discloses an ultra-high temperature hot forging method in which a steel material is heated and forged at a temperature higher than a conventional hot forging temperature of 1150 to 1250 ° C. While the conventional hot forging is set to 1150 to 1250 ° C., the heating temperature in the ultra-high temperature hot forging described in Patent Document 2 is such that the lower limit is 45 ° C. below the solidus temperature and the upper limit is the liquid phase. The temperature is 20 ° C. below the line temperature. However, when the steel material is heated to that extent, the austenite grains become coarse, and there is a drawback that toughness cannot be secured if the steel material is allowed to cool after forging. Therefore, in patent document 2, surface layer rapid cooling is performed immediately after ultra high temperature hot forging.
[0003]
[Problems to be solved by the invention]
The present invention is a hot forged product that requires high strength and high toughness, and is manufactured at a high temperature hot for producing parts with a more complex shape than conventional hot forged parts and improving material yield during production. In forging, problems such as deterioration of toughness after forging are solved without quenching and tempering.
[0004]
[Means for Solving the Problems]
(1) In mass%,
C: 0.1 to 0.6%, Si: 0.2 to 2.0%,
Mn: 0.5 to 2.5%, S: 0.02 to 0.10%,
Cr: 0.1 to 1.0%, V: 0.03 to 0.3%,
Al: 0.002-0.06%, N: 0.003-0.02%
The balance is composed of Fe and unavoidable impurities, the structure is composed of ferrite and pearlite, the ferrite fraction is 0.1 to 0.6, the average ferrite grain size is 10 to 30 μm, and the total decarburized layer depth specified by JIS G 0588 An ultra-high temperature hot forged non-heat treated part characterized in that DM-T is 0.02 to 0.08 mm,
(2) By mass%, in addition to the ultra-high temperature hot forged non-tempered part described in (1),
Ti: 0.003-0.05%, Mg: 0.0002-0.005%,
Zr: 0.0002 to 0.005%
Ultra high temperature hot forged non-heat treated parts characterized by containing one or more of
(3) Ultra high temperature hot forged non-heat treated parts according to (1) or (2), wherein the tensile strength is 800 to 1000 MPa,
(4) A method for producing the non-tempered part according to any one of (1) to (3), wherein the steel material comprising the component according to (1) or (2) is in an equilibrium state with a minimum temperature. The solidus temperature in the figure × 0.94 or 1250 ° C., whichever is higher, heated to a range where the upper limit temperature is the solidus temperature, and after ultra-high temperature hot forging in the above temperature range, A method for producing an ultra-high temperature hot forged non-heat treated part, characterized in that the processed product is subjected to finish forging at a logarithmic strain of 0.2 or higher at a temperature of 900 ° C. or higher and 1250 ° C. or lower and then allowed to cool.
(5) A method for producing a non-tempered part according to any one of (1) to (3), wherein the steel material comprising the component according to (1) or (2) is at or above the solidus temperature , Heating to the heating limit region below the upper heating limit temperature Tu [° C.], setting the holding time in the heating limit region to 1 minute or less, and the lower limit temperature is the higher of the solidus temperature × 0.94 or 1250 ° C. Then, after ultra-high temperature hot forging in a range where the upper limit temperature is equal to or lower than the heating upper limit temperature Tu [° C.], the processed product is further set to 900 ° C. or higher and the upper limit temperature Tf [° C.] or lower to a logarithmic strain of 0. A method for producing an ultra-high temperature hot forged non-tempered part characterized by cooling after adding two or more finish forgings.
However, Tu = 1536- (317.0 * C% + 9.4 * Si% + 5.2 * Mn% + 140.3 * S%)
Tf = 1456- (440.0 * C% + 12.6 * Si% + 12.8 * Mn% + 35.3 * S%)
C%, Si%, Mn%, and S% are the contents of C, Si, Mn, and S expressed in mass%.
(6) The ultra-high temperature hot forging process according to (4) or (5), characterized in that ultra-high temperature hot forging is performed so that 90% or more of the material surface is in contact with the mold in ultra-high temperature hot forging Manufacturing method for tempered parts.
(7) After performing ultra-high temperature hot forging, hold at a load of 10 to 80% of the maximum load during forging until the material surface temperature is 900 ° C. or higher and 1200 ° C. or lower at the bottom dead center of the forging machine The manufacturing method of the ultra high temperature hot forging non-heat-treated part of any one of Claims (4)-(6) characterized by the above-mentioned.
[0005]
DETAILED DESCRIPTION OF THE INVENTION
The technical idea forming the basis of the present invention is as follows.
[0006]
In the present invention, hot forged non-tempered steel whose components are adjusted so that predetermined strength and toughness can be obtained while being allowed to cool as it is is used. If it is forged, the required toughness cannot be obtained. However, hot forging is usually formed not only in a single process but also in a predetermined shape through several forging processes. Therefore, in the present invention, rough forming is performed by ultra-high temperature hot forging, and subsequent finish forming is performed at a temperature of 900 ° C. to 1250 ° C. That is, in order to form a more complicated shape than the conventional hot forging, it is formed into a substantially product shape by rough forming of ultra high temperature hot forging. And since high toughness cannot be ensured as it is, toughness is ensured by performing finish molding at a predetermined temperature and strain. In addition, since the surface ferrite decarburized layer after forging is in a predetermined range, the surface ferrite is decarburized even when heated to near the solidus, thus preventing melting cracks on the surface, particularly at the corners. Is possible.
[0007]
The present invention is described in detail below.
[0008]
C is an element effective for strengthening steel, but if it is less than 0.1%, sufficient strength cannot be obtained. On the other hand, if added excessively, toughness decreases, so the upper limit of the amount added is 0.6%.
[0009]
Si acts as a deoxidizer and is used as a solid solution strengthening element. If it is less than 0.2%, the action as a deoxidizing material is insufficient, and if added excessively, the strength is increased more than necessary and the toughness is lowered, so the upper limit of the addition amount is 2.0%.
[0010]
Mn performs strength adjustment and deoxidation action. If it is less than 0.5%, the strength is insufficient, and if it exceeds 2.5%, the toughness is lowered and cracking occurs during hot rolling, making it difficult to produce.
[0011]
S is an element indispensable for improving the machinability, and the effect is expected with more S than 0.02%. However, if it exceeds 0.10%, the toughness is lowered.
[0012]
Cr, like Mn, is an element for supplementing strength, and it is necessary to add 0.1% or more to obtain the effect. To supplement the strength, it can be added up to 1.0%, but if it exceeds 1.0%, the toughness deteriorates.
[0013]
V improves toughness by solid solution strengthening and precipitation strengthening. To obtain this effect, 0.03% or more must be added. However, even if it is added excessively, the improvement of the effect is small, but rather the toughness is lowered, so the upper limit was made 0.30%.
[0014]
Al is an element effective for deoxidation of steel and refinement of crystal grains, but if less than 0.002%, there is no effect. On the other hand, if added excessively, the upper limit of the addition amount is set to 0.06% in order to reduce toughness.
[0015]
N is an element necessary for the precipitation strengthening by generating V carbonitride, but if it is less than 0.003%, a sufficient effect cannot be obtained. On the other hand, if added excessively, the toughness deteriorates due to the dissolved N, so the upper limit of the amount added is 0.02%.
[0016]
Ti produces nitrides and carbides. Since nitride remains without dissolving at high temperatures, it is effective in preventing austenite coarsening during heating. In addition, carbides are finely dispersed and effective for precipitation strengthening. If the content is less than 0.003%, these effects do not appear. If the content exceeds 0.05%, the toughness deteriorates, so the lower limit of the amount added is 0.003% and the upper limit is 0.05%.
[0017]
Mg and Zr are both elements that form an oxide, sulfide, or a composite thereof, and have an effect of suppressing austenite coarsening during heating. Further, since these oxides become MnS precipitation nuclei, machinability is also improved. In any case, if less than 0.0002%, there is no effect, and if over 0.005%, the toughness deteriorates, so the upper limit of the addition amount is made 0.005%.
[0018]
The structure can be classified into ferrite + pearlite, bainite, and martensite depending on the cooling rate after forging. However, in the present invention, the structure of ferrite + pearlite is used as a structure that can obtain high strength and high toughness while being allowed to cool after forging.
[0019]
As the ferrite fraction increases, the toughness improves, but a ferrite fraction exceeding 0.6 causes a decrease in strength. Moreover, since the toughness is low when the ferrite fraction is low, the fraction is set to 0.1 to 0.6. The ferrite fraction can be controlled by appropriately adding the C amount and Si amount of the steel material. The ferrite fraction is defined as a value obtained by observing 3 to 5 fields of view at 200 to 500 times with an optical microscope, binarizing the contrast between ferrite and pearlite and calculating the fraction with an analysis system.
[0020]
The average grain size of ferrite can be reduced by lowering the finishing temperature and carrying out strong processing, but this results in a rapid increase in mold load, and the target value of grain size is more than necessary to improve mechanical properties. If it is made small, forging itself becomes difficult. Therefore, the lower limit of the average ferrite grain size is set to 10 μm in consideration of the level that can be actually manufactured. Further, if the average ferrite particle diameter exceeds 30 μm, toughness deterioration is caused, so the average ferrite particle diameter is set to 10 to 30 μm. In order to control the average ferrite grain size, addition of Ti, Mg, Zr, etc. is important as well as optimization of the amount of Al added to the steel material. Moreover, not only steel components but also finishing forging after ultra-high temperature forging is necessary, and control of forging temperature and processing amount is important. In particular, in order to reduce the average ferrite grain size, it is necessary to carry out processing with a logarithmic strain of 0.2 or more in the temperature range of 900 to 1250 ° C., as will be described later, as finishing forging conditions. The average ferrite particle diameter is defined as a value obtained by observing 3 to 5 fields of view at 200 to 500 times with an optical microscope and obtained by a cutting method.
[0021]
Further, the total decarburized layer depth DM-T after forging, but a decarburized layer exceeding 0.08 mm significantly reduces the strength of the forged product. In particular, in a hot forged product having a shaft portion, the shaft portion is often subjected to torsional deformation, and the surface strength is important. Decarburization occurs in the heating stage of ultra-high temperature hot forging, but in high-speed heating where the total decarburized layer depth is less than 0.02 mm, melt cracking occurs at the surface of the material, especially at the corners where heat is concentrated during heating. There is a problem that occurs. The total decarburized layer depth is 0.08 mm or less, depending on the heating device that requires rapid heating with high frequency, etc., but the feed rate of the material that controls the frequency etc. or moves through the heating device It is possible to control the total decarburization layer depth by controlling the material heating time such as controlling.
[0022]
The lower limit of the tensile strength was limited to 800 MPa in terms of weight reduction of the forged product. On the other hand, when it exceeds 1000 MPa, the toughness is remarkably lowered, and the cutting life and the die life are also remarkably lowered. Therefore, the upper limit is set to 1000 MPa or less.
[0023]
Next, a manufacturing method will be described.
[0024]
In order to reduce the deformation resistance by ultra-high temperature hot forging, the heating temperature was set to the lower limit of the solidus temperature × 0.94 or 1250 ° C., whichever is higher. This is because the deformation resistance is not sufficiently lowered at a temperature lower than this, and the material flow is not sufficiently performed. The reason why the upper limit temperature is the solidus temperature is that crystal grain boundaries and the like melt at a temperature exceeding the solidus. When the inside of the raw material is melted, the melted portion is easily vacated at the time of internal forging. Therefore, forging under a condition capable of reducing the vacancies is essential. Therefore, in the present invention, the upper limit temperature is the solidus temperature.
[0025]
The solidus temperature is estimated by a unidirectional solidification experiment used to observe the solidification process of precipitates described in Iron and Steel 73, Vol. 4, No. S196, 1987. A furnace is provided with a temperature gradient using high-frequency heating and a carbon susceptor, the rod is heated in the furnace, and then rapidly cooled. The solidus temperature of the material was estimated by correlating the temperature at each position of the bar with the structure from observation of the internal structure of the rapidly cooled bar. The details of the unidirectional solidification experiment are Matsumiya (Matsumiya) and four others, “Numerical Analysis of Segregation in Continuous Casting Slabs (Japan)”. (Transactions of the Iron and Steel Institute of Japan), Japan Iron and Steel Institute, November 1984, Vol. 24, No. 11, p. 875, 876.
[0026]
Ultra-high temperature hot forging is performed in the same temperature range as heating. At a temperature lower than the solidus temperature × 0.94 or 1250 ° C., whichever is higher, the deformation resistance is high, the mold life is reduced, and the material flow is poor, so that it cannot be molded into a complex shaped part. Further, at a temperature higher than the solidus temperature, a molten portion is generated inside the material due to processing heat generated during forging, leading to toughness deterioration.
Finish forging after rough forming by ultra-high temperature hot forging. This finish forging is intended to give a product shape to a predetermined shape and to provide toughness in the present invention, because the effect of toughness is small unless a forging process having a logarithmic strain of 0.2 or more is performed. In addition, the processing rate in finish forging is preferably about 0.5 to 2.0 in terms of logarithmic strain because it is formed into an almost product shape by ultra-high temperature hot forging. The actual strain is not uniform depending on the site, but the strain distribution is not uniform, but in order to clarify the range in which excellent toughness can be ensured, it is necessary to express the degree of processing by a predetermined formula. is there. Therefore, according to the present invention, the logarithmic strain at the time of forging is calculated from the material shape before and after processing, and if it is an upsetting part, the logarithmic strain is obtained by ln (the material height before processing / the material height after processing) and The logarithmic strain is determined by ln (material cross-sectional area before processing / material cross-sectional area after processing). If the value obtained by this formula is 0.2 or more, excellent toughness can be secured.
The reason why the finish forging is performed at 1250 ° C. or less is that at higher temperatures, the austenite grains after forging are large and the toughness is low. In order to improve the toughness, it is desirable to increase the ferrite fraction by processing at a low temperature, but at a temperature lower than 900 ° C., the deformation resistance increases rapidly, leading to a decrease in mold life.
Even when the steel material is heated to a temperature higher than the solidus temperature, if the heating upper limit temperature is set to Tu [° C.] or less and the heating time is short, melting of crystal grain boundaries and the like can be prevented. The heating upper limit temperature Tu [° C.]
Tu = 1536- (317.0 * C% + 9.4 * Si% + 5.2 * Mn% + 140.3 * S%)
It can ask for. Here, C%, Si%, Mn%, and S% are the contents of C, Si, Mn, and S expressed in mass%.
If the holding time for which the steel material is exposed to the heating limit region not lower than the solidus temperature and not higher than the heating upper limit temperature Tu [° C.] is within 1 minute, ultra high temperature hot forging is performed without generating voids inside. It is possible to suppress the coarsening of the austenite grains. Since the steel material may be heated to the solidus temperature and immediately subjected to ultra-high temperature hot forging, the lower limit of the holding time may be 0 minutes.
When performing such heating, it is preferable to increase the heating rate, and high-frequency heating or current heating is preferable. Moreover, when heating and holding, it is preferable to cool only the surface layer of the steel material, such as blowing air between the high-frequency coil and the steel material, to prevent melting of the surface layer portion.
Moreover, the coarsening of the austenite grain diameter after an ultra high temperature forging process is also suppressed by preventing the coarsening of the austenite grain before an ultra high temperature forging process. Therefore, even if the upper limit of the finish forging temperature is higher than 1250 ° C., the ferrite average grain size becomes fine if the finish upper limit temperature Tf [° C.] or less, and the toughness after forging can be ensured. Finish upper limit temperature Tf [° C]
Tf = 1456- (440.0 * C% + 12.6 * Si% + 12.8 * Mn% + 35.3 * S%)
Can be obtained. Here, C%, Si%, Mn%, and S% are the contents of C, Si, Mn, and S expressed in mass%.
[0027]
One of the objects of the present invention is to improve the material yield. Therefore, the material yield can be improved by improving the material flow in ultra-high temperature hot forging, but at that time, it can be molded to a shape close to the product shape by using a sealed mold or semi-sealed mold, and the material yield is Greatly improved. In the conventional hot forging, even if a mold having a high deformation resistance and a high sealing degree is used, the material is not filled into the mold. However, since the deformation resistance is low in ultra-high temperature hot forging, even if a mold having a high degree of sealing is used, the material filling in the mold is improved. However, unless ultra-high temperature hot forging is performed so that 90% or more of the material surface is in contact with the mold, the material yield improvement effect is small and the product shape is difficult to achieve. In addition, although the material yield is improved by performing ultra-high temperature hot forging with a mold having a high hermeticity, 98% or less is desirable because the forging load increases at the same time. Furthermore, instead of using a highly sealed mold, lowering the bottom dead center of the ram of the forging machine and reducing the size of the material to be used can reduce the contact rate and increase the material yield. .
[0028]
In order to perform finish forging at a predetermined temperature, it may take time for the material temperature to drop after ultra-high temperature hot forging. In this case, the accuracy of the product shape deteriorates due to thermal strain until the finish forging. Therefore, product accuracy is improved by freezing the shape after performing ultra-high temperature hot forging. Hold at the bottom dead center of the forging machine until the surface temperature of the material reaches 900-1200 ° C. The reason why the load is held until the surface temperature becomes 1200 ° C. or less is to reduce thermal strain. In addition, it is desirable to hold the load until it becomes 1000 ° C. or less. The reason why the lower limit temperature is set to 900 ° C. or more is that when the temperature is lower than 900 ° C., the forging load is rapidly increased in the next finish forging, and the die life is reduced. Further, the reason why the holding load is 10% or more of the maximum load at the time of forging is that if it is less than this, the suppression of thermal strain is insufficient. Note that an increase in the holding load affects the mold life, so the upper limit is 80% or less of the maximum load. Instead of holding the load at the bottom dead center of the forging machine, it is possible to cool the surface with water or the like.
[0029]
【Example】
Experiments of the present invention and comparative examples were performed using steel materials having chemical components shown in A to AA in Table 1. Steel types A to Q are target steel types of the present invention examples, and steel types R to AA are steel types used in the comparative examples. In the table, the adjacent component (P) and the solidus temperature (Ts) are also shown for reference. The solidus temperature is a temperature estimated from a unidirectional solidification test of a rod-shaped material of φ15 × 250 mm.
[0030]
Table 2 shows the results of examining the strength and toughness after ultra-high temperature hot forging-finish forging using the steel types shown in Table 1. In the experiment, a φ60 × 60 mm material was heated to a predetermined temperature at a high frequency. In addition, the frequency at the time of a heating was 3-5 KHz, it heated from room temperature to 1250 degreeC at the speed | rate of 5 degree-C / sec, and was heated to predetermined temperature at 1 degree-C / second after that. After reaching a predetermined temperature, it was held for about 30 seconds and then subjected to forging. Since the sample is moved to the forging machine after the high frequency heating, the material temperature is lowered. Accordingly, Table 2 shows the heating temperature as Tk (° C.) and the temperature immediately before forging as Tt (° C.). The temperature Tt immediately before forging is the result of measuring the temperature of the material surface with a radiation thermometer. Forging was performed by a compression tester having a hydraulic servo mechanism at a ram speed of 200 mm / s. In the ultra-high temperature hot forging, a φ60 × 60 mm sample was placed horizontally, and compression was performed at a compression rate of about 50% using a flat platen. By compressing a 60 mm-high material to 30 mm before processing, the compression ratio becomes 50%, and the logarithmic strain is calculated as ln (material height before processing / material height after processing). Becomes 0.7 (= ln (60/30)). Further, finish forging is further performed after ultra high temperature forging. When a material having a height of 30 mm is further processed to 16 mm by ultra high temperature forging, the logarithmic strain in finish forging is 0.6 (= ln (30/16)). It becomes.
[0031]
It takes time for the hydraulic pressure of the testing machine to recover after ultra high temperature forging and finish forging, and since it is lower than the temperature during ultra high temperature hot forging, the surface temperature of the material immediately before finishing is measured with a radiation thermometer. The heated results are shown in Table 2. After finishing forging, after cooling, the ferrite fraction, the average ferrite grain size, and the total decarburized layer depth were measured. The decarburization layer depth was 200-500 times with an optical microscope, and the decarburization layer depth was measured at 10 locations using an eyepiece with reading dimensions, and the average value of 8 locations excluding the maximum and minimum was taken. In addition, tensile test pieces and Charpy impact test pieces were collected from the compressed samples. Table 2 shows the tensile strength and impact value. The impact value is the absorbed energy at -50 ° C.
[0032]
From Table 2, Examples 1 to 17 of the present invention have a tensile strength of about 800 to 900 MPa and an impact value of 20 J / cm. 2 That's it. On the other hand, Comparative Examples 1 to 10 are cases where the steel types of Comparative Examples R to AA in Table 1 are used, and the impact value is 20 J / cm in all cases. 2 I have n’t been to In particular, in Comparative Example 5, the strength is as low as 787 MPa. In Comparative Example 11, since the temperature before finish forging is as low as 850 ° C., the ferrite fraction is higher than the predetermined value and the strength is low. In Comparative Example 12, the pre-finishing temperature is lower, and when finishing forging at such a low temperature, the ferrite grain size becomes smaller than a predetermined value, and high strength can be obtained as compared with the present invention. Is about 3000 KN, which is four times or more the load of about 700 KN of the present invention, and the load on the mold and forging machine increases rapidly. Although the comparative example 14 is a case where the process at the time of finish forging is small, all had low toughness after forging. Comparative Example 15 was a case where the heating temperature at the time of ultra-high temperature hot forging and the temperature immediately before forging were high. The decarburized layer was deepened and did not reach the predetermined strength. In Comparative Example 15, a part of the sample was melted during heating. In Comparative Example 16, the holding temperature time after heating was increased to 5 minutes, and in this case as well, the decarburization increase was deep and the strength was low.
[0033]
Table 3 shows a comparison when various processing conditions are further changed. In the example of the present invention, the load at the time of ultra-high temperature hot forging is 282 to 416 KN, and at the time of finish forging is 701 to 796 KN. On the other hand, Comparative Example 17 is a case where the heating temperature is low, and Comparative Example 18 is a case where the temperature immediately before the ultra-high temperature hot forging is made lower than the predetermined temperature by increasing the time until the ultra-high temperature hot forging although it is heated to the predetermined temperature. However, both are high loads of 600KN or more. In Comparative Example 19, the temperature immediately before the finish forging is low, but the finish forging load at this time is 1826 KN, which is more than double that of the finish forging in the present invention example. A high forging load will lead to a decrease in mold life. Although the comparative example 20 is a case where finish forging is not performed, it turns out that a toughness value is low. Comparative Example 21 is a case in which ultra-high temperature hot forging was performed immediately without holding for 30 seconds after heating to a predetermined temperature in order to reduce the depth of the decarburized layer. The temperature is measured on the surface of the sample. Although the surface temperature has reached the specified temperature, the temperature is not maintained, so the center material is not heated to the specified temperature. The load of 864 KN was higher than the load at the time of ultra-high temperature hot forging in the present invention.
[0034]
Table 4 shows a case where a predetermined load is held by a ram of the compressor on the sample after ultra-high temperature hot forging in order to confirm the effect of the invention according to (7). In any case, ten samples were forged under predetermined conditions, and the height of the forged product after finish forging was measured. The difference in height between the object with the maximum height and the object with the minimum height among the 10 is indicated by Hb in the same table. Note that the heating temperature, pre-forging temperature, load, and the like are the results of measurement on one of 10 samples. In the example of the present invention, Hb was 0.24 to 0.34 mm. On the other hand, Reference Example 1 is a case where the holding load is small and Reference Example 2 is a case where the holding temperature is high. There were many variations. Reference Example 3 is a case where the holding temperature is low, and although the Hb is small, if the holding temperature is lowered, the temperature before forging at the time of finish forging also becomes low, and the load at the time of finish forging becomes nearly doubled, thereby deteriorating the die life. I will let you.
[0035]
Table 5 shows that the forging was performed with the component shape as shown in FIG. 1 in order to examine the forging yield in order to confirm the effect of the invention according to (6). The forging machine used a general crank press, and heating was performed in a high frequency furnace. Using a round bar of the size shown in the table, after forging at ultra-high temperature, finish forging was performed, and trimming into the shape of FIG. By weighing the forged product after trimming, the material yield was obtained from the weight ratio with the material before forging. The contact ratio between the material and the mold in the ultra-high temperature hot forging was adjusted by moving the bottom dead center of the forging machine ram and the mold up and down. Tensile test pieces and impact test pieces were taken from the forging machine and the arm of the forged product. Although it is the processing amount of the arm part, the arm part processed the raw material of diameter 53-58mm to height 33mm in ultra high temperature hot forging. Then, it is 18 mm by finish forging, and the logarithmic strain in finish forging is 0.6 (= ln (33/18)). The impact value was represented by a value at −50 ° C. First, as an experiment, an experiment of Comparative Example 22 corresponding to normal hot forging was performed. By using a φ58 × 177 mm material at a contact rate of 92%, it was possible to process so as not to generate an unformed part, and the material yield at that time was 62%. Next, the present invention example 24 and the present invention example 25 were subjected to experiments in ultra-high temperature hot forging. Invention Example 24 was a sample having a size smaller than that of Comparative Example 22, and could be processed without an unmolded portion even with a contact rate of 93%. This is because the hot forging of Comparative Example 22 has a low heating temperature and therefore the material flow is poor, and a contact rate of 92% cannot be secured unless sufficient flash is generated from the mold flash part. On the other hand, in Invention Example 24, since the material flow was good, the contact rate could be ensured without producing burrs. As a result, the yield in Invention Example 24 with a small amount of burrs was remarkably improved. Invention Example 25 is a case where the bottom dead center of the forging machine ram is adjusted to improve the contact rate, but the yield is also improved to 87%. Reference Example 4 and Reference Example 5 are cases where the bottom dead center was adjusted to a contact rate of 85%, but in Reference Example 4 forged with a sample of the same size as Example 24 of the present invention, unprocessed parts were generated. . Therefore, it is necessary to use a material having a diameter of 55 × 174 mm in order to prevent the unprocessed portion from occurring. This is Reference Example 5, but the material yield was as low as 71%. From the above, it can be seen that the present invention can be forged at a high yield, and can be manufactured at a low manufacturing cost.
Table 6 shows that in order to confirm the effect of the invention according to the above (5), the steel material is heated within the solid line temperature to Tu [° C.] or less within 1 minute, and the upper limit temperature of finish forging is set to Tf [° C.] or less. This is a comparison. Using samples of φ60 × 60 mm composed of the components of steels A to D in Table 1, about 50% ultra-high temperature hot forging was performed under the conditions shown in Table 6, then finishing forging and allowing to cool. In addition, a sample with a thermocouple attached to the center is prepared, and ultra-high temperature hot forging and finish forging are separately performed under the conditions shown in Table 6, and the center temperatures before ultra-high temperature hot forging and before finish forging are set. It was measured. The results are also shown in Table 6.
In Table 6, Examples 26 to 33 of the present invention can be forged without causing any melt cracking, and the impact value at −50 ° C. after forging is also 20 J / cm. 2 That's it. On the other hand, Comparative Examples 23 and 24 had a high heating temperature, melted during heating, and could not perform ultra-high temperature hot forging. Therefore, “−” is entered in the columns of the temperature before forging, forging load, finish forging, and forging after forging in the table. In Comparative Examples 25 and 26, the heating time was longer than the solidus temperature, and forging cracks occurred although the ultra-high temperature hot forging load was low. Therefore, finish forging could not be performed, and “-” was entered in the column of the product after finish forging and forging in the table. Since Comparative Examples 27 and 28 have a high finish forging temperature, and Comparative Examples 29 and 30 have a small distortion in finish forging, the impact value at −50 ° C. after forging is 20 J / cm. 2 Not reached.
The reference example is the result of an experiment under the condition (4). In the reference examples 6 to 9, the ultra high temperature hot forging load is 382 to 477 kN and the finishing forging load is 1826 to 2225 kN, whereas in the inventive examples 26, 27, 32, and 33 which are equivalent strain conditions, the high forging load is extremely high. A warm forging load is 211 to 285 kN, and a finishing forging load is 1495 to 1539 kN, and a significant load reduction effect is recognized.
[0036]
[Table 1]
[Table 2]
[Table 3]
[Table 4]
[Table 5]
[Table 6]
【The invention's effect】
According to the present invention, the formability of the steel material is remarkably enhanced, so that it is possible to process a complex shaped part that could not be achieved in the past and to process with a high material yield. This realizes the weight reduction of the parts and the production with higher productivity and lower cost than the conventional one, which brings about a great effect in the production of the machine parts.
[Brief description of the drawings]
FIG. 1 is a diagram showing a forged product shape and schematic dimensions in a forging experiment.
Claims (7)
Ｃ ：０．１〜０．６％
Ｓｉ：０．２〜２．０％
Ｍｎ：０．５〜２．５％
Ｓ ：０．０２〜 ０．１０％
Ｃｒ：０．１〜１．０％
Ｖ ：０．０３〜０．３％
Ａｌ：０．００２〜０．０６％
Ｎ ：０．００３〜０．０２％
を含有し、残部Ｆｅおよび不可避不純物からなり、組織はフェライトとパーライトからなり、フェライト分率０．１〜０．６、フェライト平均粒径１０〜３０μｍ、ＪＩＳ Ｇ ０５８８で 規定する全脱炭層深さＤＭ−Ｔが０．０２〜０．０８ｍｍであることを特徴とする超高温熱間鍛造非調質部品。% By mass
C: 0.1 to 0.6%
Si: 0.2-2.0%
Mn: 0.5 to 2.5%
S: 0.02-0.10%
Cr: 0.1 to 1.0%
V: 0.03-0.3%
Al: 0.002 to 0.06%
N: 0.003-0.02%
The balance is composed of Fe and unavoidable impurities, the structure is composed of ferrite and pearlite, the ferrite fraction is 0.1 to 0.6, the average ferrite grain size is 10 to 30 μm, and the total decarburized layer depth specified by JIS G 0588 An ultra-high temperature hot forged non-heat treated part characterized in that DM-T is 0.02 to 0.08 mm.
Ｔｉ：０．００３〜０．０５％
Ｍｇ：０．０００２〜０．００５％
Ｚｒ：０．０００２〜０．００５％
の中から１種または２種以上を含有することを特徴とする超高温熱間鍛造非調質部品。In addition to the ultra-high temperature hot forged non-tempered part according to claim 1, in addition to Ti: 0.003 to 0.05% by mass%
Mg: 0.0002 to 0.005%
Zr: 0.0002 to 0.005%
An ultra-high temperature hot forged non-heat treated part characterized by containing one or more of the above.
ただし、Ｔｕ＝１５３６−（３１７．０×Ｃ％＋９．４×Ｓｉ％＋５．２×Ｍｎ％＋１４０．３×Ｓ％）
Ｔｆ＝１４５６−（４４０．０×Ｃ％＋１２．６×Ｓｉ％＋１２．８×Ｍｎ％＋３５．３×Ｓ％）
であり、Ｃ％、Ｓｉ％、Ｍｎ％、Ｓ％は、質量％で表したＣ、Ｓｉ、Ｍｎ、Ｓの含有量である。It is a method of manufacturing the non-tempered part of any one of Claims 1-3, Comprising: The steel material which consists of a component of Claim 1 or 2 is more than solidus temperature, and heating upper limit temperature Tu [degreeC The following heating limit region is heated, the holding time in the heating limit region is set to 1 minute or less, the lower limit temperature is set to the higher of the solidus temperature × 0.94 or 1250 ° C., and the upper limit temperature is set to the upper limit of heating. After ultra-high temperature hot forging in the range of temperature Tu [° C.] or less, the workpiece is further subjected to finish forging at a temperature of 900 ° C. or higher and a finish upper limit temperature Tf [° C.] or lower with a logarithmic strain of 0.2 or higher. A method for producing an ultra-high temperature hot forged non-tempered part characterized by cooling after standing.
However, Tu = 1536- (317.0 * C% + 9.4 * Si% + 5.2 * Mn% + 140.3 * S%)
Tf = 1456- (440.0 * C% + 12.6 * Si% + 12.8 * Mn% + 35.3 * S%)
C%, Si%, Mn%, and S% are the contents of C, Si, Mn, and S expressed in mass%.
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