JP4088546B2

JP4088546B2 - Manufacturing method of aluminum alloy forging with excellent high temperature characteristics

Info

Publication number: JP4088546B2
Application number: JP2003090660A
Authority: JP
Inventors: 廉樹上高原; 俊弘桂; 学中井; 泰彰渡辺
Original assignee: Kobe Steel Ltd
Current assignee: Kobe Steel Ltd
Priority date: 2002-03-29
Filing date: 2003-03-28
Publication date: 2008-05-21
Anticipated expiration: 2023-03-28
Also published as: JP2004002987A

Description

【０００１】
【発明の属する技術分野】
本発明は、2000系アルミニウム合金鍛造材 (以下、アルミニウムを単にAlと言う) に関し、高温特性 (耐熱性および高温耐力) に優れたAl合金鍛造材の製造方法に関するものである。
【０００２】
【従来の技術】
ロケットや航空機などの航空・宇宙機材用、鉄道車両、自動車、船舶などの輸送機材用、あるいはエンジン部品、コンプレッサーなどの機械部品用、具体的には、回転ローターや回転インペラー或いはピストンなどの、特に100 ℃を超える高温の使用環境となるAl合金製部品には、高温特性に優れたAl合金鍛造材が用いられる。この高温特性とは、前記高温下での耐クリープ特性および高温耐力である。
【０００３】
従来、これらの所謂耐熱性Al合金鍛造材には、AA規格乃至JIS 規格の 2000 系( 以下、単に2000系と言う)Al 合金が用いられている。この種Al合金としては、2219、2618などがある。しかし、これらの2000系 Al 合金は、120 ℃を越える高温では、長時間使用すると強度の低下が著しい。
【０００４】
このため、120 ℃を越える高温使用環境でのクリープ特性や高温耐力を改善するために、近年では、2219Al合金にMgを0.3%添加した2519Al合金(Al-6.1Cu-0.3Mn-0.15Zr-0.1V)が開発されている。また、この2519Al合金にAgを添加した2519(Ag)Al合金も開発されている。そして、これら2519Al合金や2519(Ag)Al合金に関連したAl合金も多数提案されている (例えば、特許文献１、２参照）。また、本発明者らも、高い高温特性を再現性良く保証することが可能な耐熱Al合金材を提案した。この内容は、Cu:1.5〜7.0%、Mg:0.01 〜2.0%を含み、更に、選択的にAg:0.05 〜0.7%を含む耐熱Al合金の、θ' 相および/ またはΩ相について、θ' 相の平均サイズを120 nm以下およびθ' 相の析出物間の平均間隔を100 nm以下とすること、Ω相の平均サイズを100 nm以下およびΩ相の析出物間の平均間隔を150 nm以下とすることである (特許文献３、非特許文献１参照) 。
【０００５】
【特許文献１】
特開昭62-112748 号公報
【特許文献２】
米国特許第4610733 号明細書
【特許文献３】
特開平11-302764 号公報
【非特許文献１】
軽金属学会第93回秋期大会講演概要(1997 年 10 月20日発行、233 〜 234 頁)
【０００６】
また、前記高温特性が要求される用途部品は、基本的に肉厚の円筒形状や多数の羽根を周囲に設けた複雑形状を有している。このため、Al合金材によりこれらの部品を製造する場合には、Al合金のバルク状 (塊状) の鋳塊を熱間鍛造加工（熱間鍛造後冷間鍛造することも含む）した鍛造材から切削加工により部品とされている。そして、これら用途部品は、狭い空間乃至クリアランスを高速で摺動乃至回転するため、高い寸法精度や平滑性が厳しく要求される。このため、これら用途に使用されるAl合金材には、前記高温特性に加えて高い精密切削加工性、即ち被削性が要求される。
【０００７】
このため、本発明者らは、高速動部品用の耐熱Al合金鍛造材の高い高温特性とともに高速動部品への切削加工における被削性を保証するために、Al合金鍛造材の溶体化処理後のミクロ組織がθ' 相および/ またはΩ相を有するとともに、結晶粒径を500 μm 以下の等軸再結晶粒とすることも提案した (特許文献４参照) 。
【０００８】
【特許文献４】
特開2000-119786 号公報
【０００９】
しかし、これらの技術により、高温特性に優れたAl合金鍛造材を冶金的に設計したとしても、実際に製造されるAl合金鍛造材において、溶体化処理および焼入れ処理後の高温の人工時効硬化処理を施しても、耐力が向上せず、この種Al合金鍛造材 (耐熱Al合金鍛造材) に要求される人工時効硬化処理後の耐力が低くなり、高温使用時の耐力も低くなる場合が生じる。このため、本発明者らは、溶体化処理後の焼入れ速度の影響に注目し、400 ℃から290 ℃の間の平均冷却速度が30000 ℃/ 分以下と焼入れ速度 (冷却速度) が遅く (小さく) なる場合には、特に、Al合金鍛造材中のZr、Cr、Mnを、Zr:0.09%以下、Cr:0.05%以下、Mn:0.6% 以下に各々規制することを提案した (特許文献５参照) 。
【００１０】
【特許文献５】
特開2001-181771 号公報
【００１１】
【発明が解決しようとする課題】
しかし、高温特性に優れるように冶金的に設計されたAl合金鍛造材を、その設計高温特性通りに再現性良く製造できるためには、なお、熱間鍛造条件や溶体化焼入れ処理条件などの、実際の製造条件の改良が必要であり、開発要素が残されていた。
【００１２】
本発明はこの様な事情に着目してなされたものであって、その目的は、熱間鍛造条件や溶体化焼入れ処理条件などの製造条件を改良し、その設計高温特性通りに再現性良く製造できるAl合金鍛造材の製造方法を提供しようとするものである。
【００１３】
【課題を解決するための手段】
この目的を達成するために、本発明アルミニウム合金鍛造材の製造方法の請求項1 の要旨は、Cu:4.0〜7.0%、Mg:0.2〜0.4%、Ag:0.05 〜0.7%を含み、残部アルミニウムおよび不可避的不純物からなるアルミニウム合金鍛造材の製造方法であって、この組成からなる鋳造材を500 〜535 ℃の温度で均質化熱処理後、280 〜360 ℃の温度で熱間鍛造し、その後510 〜545 ℃の温度で溶体化および焼入れ処理し、人工時効硬化処理を施した場合のアルミニウム合金鍛造材の室温での耐力が400MPa以上である特性を有することとする。なお、合金元素含有量の% 表示は全て質量% を意味する。
【００１４】
ここにおいて、前記冶金的に設計されたAl合金鍛造材を、その設計高温特性通りに再現性良く製造するためには、熱間鍛造条件や溶体化焼入れ処理条件などが重要であることは疑いがない。しかし、本発明者らは、これら熱間鍛造や溶体化焼入れ処理の最適条件が存在することを知見した。
【００１５】
即ち、先ず、熱間鍛造の温度は、その設計高温特性通りに再現性良く製造するためには、一般的な熱間鍛造温度の条件範囲からより低温側である必要がある。また、溶体化処理温度及び溶体化処理後の焼入れ温度は、その設計高温特性通りに再現性良く製造するためには、後述する通り、本発明者らが知見した上記温度範囲にて処理することが必要である。
【００１６】
【発明の実施の形態】
以下に本発明の各要件の意義について説明する。
本発明におけるAl合金鍛造材の製造工程自体は、従来と基本的に同じである。即ち、本発明の成分範囲内に溶解調整されたAl合金溶湯を、連続鋳造圧延法、半連続鋳造法（DC鋳造法）等の通常の溶解鋳造法を適宜選択して鋳造して鋳塊を製作する。この鋳塊を500 〜535 ℃の温度で均質化熱処理後、熱間鍛造してAl合金鍛造材を製造する。なお、鍛造用の素材としては、鋳塊を押出、圧延加工した、押出材や圧延材を使用しても良い。ここにおいて、前記均質化熱処理の温度が500 ℃未満では鋳塊の晶出物が固溶せず、均質化が不十分となる。一方、前記均質化熱処理の温度が535 ℃を越えると、バーニングが生じる可能性が高くなる。したがって、前記均質化熱処理の温度は500 〜535 ℃の範囲とする。
【００１７】
ここにおいて、本発明におけるAl合金鍛造材の製造条件や製造手段は、以下に詳述する熱間鍛造温度や溶体化処理温度や溶体化後の焼入れ温度及び必要に応じて行う冷間圧縮条件などを除き、従来方法と基本的には同じである。言い換えると、Al合金鍛造材の製造条件や製造手段を大きく変えない点が、本発明の利点でもある。
【００１８】
まず、熱間鍛造の温度条件は、その設計高温特性通りに、Al合金鍛造材を再現性良く製造するために重要である。従来では、自由鍛造や型鍛造 (鍛伸鍛造) などの公知の鍛造手段を単独あるいは組み合わせて、適宜採るにせよ、Al合金鍛造材の溶体化処理後のミクロ組織を等軸結晶粒とするために、熱間鍛造温度を380 〜430 ℃程度としていた。この熱間鍛造温度が低いと、Al合金鍛造材の組織が局部的に混粒となりやすく、高温特性が低下すると認識していたためである。
【００１９】
この点、本発明では、熱間鍛造温度を再結晶温度以下の280 〜360 ℃の温度範囲とする。熱間鍛造温度が360 ℃を越えると、本発明の成分範囲内のAl合金鍛造材には粗大粒が生じやすくなる。このため、Al合金鍛造材の高温特性が低下し、高温特性に優れたAl合金鍛造材を再現性良く製造することができない。一方、熱間鍛造温度が280 ℃未満では、熱間鍛造時に割れが生じ易く、鍛造加工自体が困難となる。
【００２０】
本発明では、熱間鍛造の温度を280 〜360 ℃としても、本発明の成分範囲内のAl合金鍛造材では、溶体化および焼入れ処理の適切化により、Al合金鍛造材の調質後のミクロ組織は等軸結晶粒となり、混粒とはならない。
【００２１】
なお、Al合金鍛造材の前記ミクロ組織は、熱間鍛造の鍛練比にも影響される。したがって、Al合金鍛造材の場合に、前記ミクロ組織を等軸結晶粒とするためには、前記適宜の熱間鍛造の鍛練比を1.5 以上とすることが好ましい。鍛練比が1.5 未満であれば、Al合金鍛造材の組織が混粒となりやすい。さらに、鍛練の方向は一方向だけではなく、少なくとも、異なる2 方向で行い、各方向での鍛練比を1.5 以上とすることが更に好ましい。
【００２２】
次に、溶体化および焼入れ処理について説明する。この溶体化および焼入れ処理において、可溶性金属間化合物を再固溶し、かつ冷却中の再析出を可能な限り抑制するためには、JIS-H-4140、AMS-H-6088などに規定された条件内にて行うことが好ましい。ただし、たとえAMS-H-6088等の規格によって熱処理を行っても、溶体化処理温度が高すぎるとバーニングを生じ、機械的性質を著しく低下させる。そして、溶体化処理温度が下限以下の温度であると人工時効硬化処理後の室温での耐力が400MPa以上とならず、また溶体化自体も困難となる。従って、溶体化処理温度の上限は545 ℃とし、下限は510 ℃とする。
【００２３】
ここで、φ100 mm程度までの小物部品やピストンなどの用途において、残留応力が比較的大きくても、例えば切削などの加工上問題とならない製品については、溶体化および焼入れ処理後に人工時効硬化処理を施し、調質T6材とすることが望ましい。この場合、残留応力が比較的大きくなっても、高い強度特性及び高温特性を得る為に、焼入れ温度は40℃以下であることが望ましい。また、この焼入れ温度が高いと、人工時効硬化処理後の室温での耐力を400MPa以上とするのが困難となる。
【００２４】
一方、ロータなど大型の製品では、焼入れ処理時に、製品表面と中央部との冷却速度が大きく異なるため、製品表面には10kgf/mm²を越える高い残留応力が発生する。このような高い残留応力が発生すると、製品の切削加工時に大きな歪みが生じ、精密な切削加工が極めて困難となる。また、最悪の場合、切削加工中に残留応力による割れなどの破壊が生じることもある。例え、切削加工中に割れなどの破壊が生じなくても、材料中に残存する晶出物等の金属間化合物を起点として、あるいは製品搬送中に生じた僅かな表面傷等を起点として、製品の長期間使用中に、き裂が伝播成長しやすく、最終破断に至る可能性もある。したがって、ロータなど残留応力が問題となる製品については、残留応力を好ましくは3.0kgf/mm²以下に除去乃至低減するため、溶体化処理後の水焼入れ温度を90℃以上の比較的高温とし、その後人工時効硬化処理を施し、調質T61 材とすることが好ましい。
【００２５】
また、用途によっては、製品の大小に関わらず、残留応力が厳しく管理される製品もある。このような製品については、残留応力を極力小さくすべく、冷間圧縮乃至冷間加工を加えて、残留応力を好ましくは3kgf/mm²以下に除去乃至低減し、人工時効硬化処理を施して調質T652材とすることが好ましい。これらの製品では、残留応力を好ましくは3kgf/mm²以下に除去乃至低減し、高い強度特性及び高温特性を得る為、焼入れ温度は40℃以下であることが好ましい。この焼入れ温度が高いと、人工時効硬化処理後の室温での耐力を400MPa以上とするのが困難となる。前記冷間圧縮乃至冷間加工の冷間圧縮 (加工) 量が小さいと十分な残留応力の低減効果が得られない。一方、冷間圧縮量が大きいと、人工時効硬化処理中や高温での使用中に、θ' 相の析出量が増加する為、耐力が低下しやすい。従って、冷間圧縮 (加工) は、圧縮 (加工) 率1 〜5 % とすることが好ましい。
【００２６】
その後、これらAl合金鍛造材は前記用途部品に加工される。勿論、Al合金鍛造材を、前記用途製品に加工後に、溶体化、焼入れ処理および冷間圧縮や人工時効硬化処理などを適宜行っても良い。
【００２７】
溶体化処理および焼入れ処理などの調質 (熱処理) に用いる炉はバッチ炉、連続焼鈍炉、溶融塩浴炉、オイル炉などが適宜使用可能である。また、焼入れに際しての冷却手段も、ユーコンクウェルチャント、水浸漬、温水浸漬、沸騰水浸漬、水噴射、空気噴射などの手段が適宜選択可能となる。
【００２８】
このようにして得られた本発明Al合金鍛造材のミクロ組織は、500 μm 以下の、好ましくは10〜500 μm の範囲の、更に好ましくは50〜300 μm の範囲の、ほぼ一定サイズの微細な再結晶粒 (等軸再結晶粒) である。そして前記混粒組織に見られるような、粒径が1 μm 以下の微細な再結晶粒( 或いは亜結晶粒) が集合体化した集団や、数mm〜数cm程度の粗大な再結晶粒、あるいは残存する鋳塊組織もなく、良好なクリープ特性などの高温特性と被削性とを兼ね備える。
【００２９】
ただ、本発明における好ましい等軸再結晶粒の組織とは、前記一定サイズの等軸再結晶粒が100%のみの組織を必ずしも意味するものではなく、前記被削性やクリープ破断強度などの高温特性を低下させない範囲での、鋳造組織や混粒組織の混入は許容する。例えば、粒径が1 μm 以下の微細な再結晶粒( 或いは亜結晶粒) は、単一の結晶粒が個々に分散して存在しても、前記被削性やクリープ破断強度などの高温特性を低下させない。しかし、これがお互いにくっついた形で集団化乃至集合体化した場合に被削性や高温特性を低下させるようになる。したがって、この点からは、溶体化処理後のミクロ組織において、集合体化している1 μm 以下の微細再結晶粒の面積率は10% 以下とすることが好ましい。
【００３０】
なお、本発明で言う等軸再結晶粒の特定および混粒組織の有無は、試料を電解エッチング等によりミクロエッチングを行い、これを50〜400 倍の光学顕微鏡により観察乃至測定可能である。
【００３１】
次に、本発明のAl合金鍛造材組織において、高温耐力やクリープ破断強さなどの高温特性をより高めるためには、溶体化処理および焼入れ処理後に、160 〜190 ℃×7 〜60時間の範囲から選択することによって、Al合金の(100) 面に析出するθ' 相、(111) 面に析出するΩ相を析出させることが、好ましい。人工時効硬化処理によるこれらの析出がないと、前記人工時効硬化処理された場合でも180 ℃などの温度での高温耐力が低くなる。
【００３２】
なお、Al合金鍛造材組織中のθ' 相とΩ相の析出状態の同定は、50000 倍の透過型電子顕微鏡(TEM) により、組織を観察して行える。
【００３３】
本発明Al合金鍛造材における、化学成分組成について説明する。本発明のAl合金の化学成分組成は、基本的に2519 或いは2618などのAl合金および2519にAgを加えた2519(Ag)系Al合金の成分規格として良いが、より具体的な用途および要求特性に応じて、以下に説明する成分組成範囲から適宜選択しうる。先ず、積極的含有元素について述べる。
【００３４】
(Cu:4.0 〜7.0%)
Cuは本発明Al合金鍛造材の基本成分であり、固溶強化及び析出強化の双方の作用により、主としてAl合金鍛造材の本発明用途において要求される、常温と高温のクリープ特性および高温耐力を確保するために必須である。より具体的には、Cuは、前記した通り、高温の人工時効硬化処理時に、θ' 相やΩ相を、Al合金の(100) 面や(111) 面に微細でかつ高密度に析出させ、人工時効硬化処理後のAl合金鍛造材の強度を向上させる。この効果は4.0%以上で発揮され、Cuの含有量が4.0%未満では上述の効果が小さく、Al合金鍛造材の常温と高温での十分なクリープ特性および高温耐力が得られない。一方、Cuの含有量が7.0%を越えると、強度が高くなりすぎ、Al合金鍛造材の鍛造性が低下する。したがって、Cuの含有量は4.0 〜7.0%の範囲とする。
【００３５】
(Mg:0.2 〜0.4%)
MgもCuと同様に、固溶強化及び析出強化の双方の作用により、主としてAl合金鍛造材の常温と高温での十分なクリープ特性および高温耐力を確保するために必須である。より具体的には、MgもCuと同様に、高温の人工時効硬化処理時に、θ' 相やΩ相を、Al合金鍛造材の(100) 面や(111) 面に微細でかつ高密度に析出させ、人工時効硬化処理後のAl合金鍛造材の強度を向上させる。この効果は0.2%以上で発揮され、Mgの含有量が0.21% 未満ではこの効果が発揮されず、Al合金鍛造材の常温と高温での十分なクリープ特性および高温耐力が得られない。一方、Mgの含有量が0.4%を越えると、強度が高くなりすぎ、溶体化処理時にバーニングと称される割れが発生したり、鍛造性を低下させる可能性が高くなる。したがって、Mgの含有量は0.2 〜0.4%の範囲とする。
【００３６】
(Ag:0.05〜0.7%)
AgはAl合金鍛造材中において、微細で均一なΩ相を形成するとともに、析出物相が存在しない領域(PFZ;solute-depleted precipitate free zone) の幅を極めて狭くすることによりAl合金鍛造材の常温および高温強度を向上させるために必須である。Agの含有量が0.05% 未満ではこの効果がなく、また一方でAgの含有量が0.7%を越えて含有しても効果は飽和する。したがって、Agの含有量は0.05〜0.7%の範囲とする。
【００３７】
(V:0.15%以下)
V は、Zr、Cr、Mnと同様に、均質化加熱処理時に、Al合金鍛造材組織中で熱的に安定な化合物であるAl-V系分散粒子を析出させ、この分散粒子が再結晶後の粒界移動を妨げる作用があるため、結晶粒の粗大化防止の効果がある。この結果、Al合金鍛造材のミクロ組織を繊維組織化して、常温強度および高温強度を向上させる効果がある。そして、安定相を粗大に析出させる作用がZr、Cr、Mnに比して比較的小さい。したがって、Al合金鍛造材の被削性や高温特性の確保をより確実に保証する目的で、結晶粒径を500 μm 以下に微細化させるために、本発明の好ましい態様では、V を0.15% 以下を選択的に含有させることが好ましい。V 含有量が0.05% 未満ではこれらの効果が得られにくく、一方、V:0.15% を越えると、溶解鋳造時に粗大な不溶性金属間化合物を生成しやすく、成形不良および破壊の原因となる。したがって、V は0.15% 以下、好ましくは0.05% 〜0.15% の範囲で含有させる。
【００３８】
以下に、規制することが好ましい元素について説明する。
Zr 、Cr、Mnは、前記V と同様に、均質化加熱処理時にそれぞれAl合金鍛造材組織中で熱的に安定な化合物であるAl-Zr 系、Al-Cr 系、Al-Mn 系の分散粒子を析出させる。そして、この分散粒子が、Al合金鍛造材のミクロ組織を繊維組織化して、常温強度および高温強度を向上させる効果を有する。
【００３９】
しかし、溶体化処理後の焼入れ処理において、400 ℃から290 ℃の間の平均冷却速度が30000 ℃/ 分以下に遅くなった場合、これらZr、Cr、Mnを含有していると、溶体化処理後の焼入れ処理において、焼入れの過程で、AlCu₂などの安定相が、前記Al-Cr 系、Al-Zr 系、Al-Mn 系の分散粒子の周囲に粗大に析出してしまう。この結果、次に高温の前記人工時効硬化処理を行っても、120 ℃の温度で100 時間使用された後に310MPa以上などの高温での耐力が得られない。したがって、Al合金鍛造材の焼入れ感受性を下げるために、好ましくは、Zr:0.09%以下、Cr:0.05%以下、Mn:0.8% 以下、各々に規制することが好ましい。
【００４０】
Feは0.15% 以下に規制することが好ましい。ただ、スクラップ等からの混入もあり、Al合金鍛造材の高温特性を向上させる効果もあるので、0.15% の含有までは許容する。0.15% を越えて含有すると、不溶性金属間化合物を生成し、成形不良および破壊の原因となりやすい。
【００４１】
SiはMgと結合してAl合金鍛造材組織中にMg₂Si およびAl-Fe-Si系の晶出物を形成する。このため、高温の人工時効硬化処理時にθ' 相やΩ相を析出させ、人工時効硬化処理後のAl合金鍛造材の強度を向上させるために必要なMgが消費されてしまうので、人工時効硬化処理後のAl合金材の強度が低下する。Mgの含有量はCuに比して、元々少ないので、このSiによる影響は大きい。また、溶体化処理により、前記大部分の晶出物は固溶するが、過剰なMg₂Si が形成されると溶体化処理においても残存して破断の起点になるため、成形性が低下する。したがって、Siは0.1%以下に規制することが好ましい。
【００４２】
この他、Tiは、結晶粒を微細化するが、過剰に添加すると粗大な金属間化合物を形成し成形加工時の破断の起点になるため、成形性が低下する。したがって、Tiは0.1%以下までの含有は許容される。
【００４３】
したがって、本発明の好ましい態様では、Al合金鍛造材の人工時効硬化処理後の耐力が低くなり、高温使用時の耐力も低くなることを防止するために、Al合金鍛造材合金中の以下の元素を、Si:0.1% 以下、Fe:0.15%以下、Zr:0.09%以下、Cr:0.05%以下、Mn:0.8% 以下、Ti: 0.1%以下に各々規制することが好ましい。
【００４４】
また、上記以外のZn、Ni、B などの元素については、本発明に係るAl合金鍛造材の高温特性やその他の特性を阻害しない範囲での含有あるいは2000系Al合金の上限規格程度は許容される。
【００４５】
【実施例】
次に本発明の実施例を説明する。表1 に示す、A 〜G の本発明範囲内の化学成分組成、H 〜L の本発明範囲外の化学成分組成のAl合金鋳塊 (500mm φ×2000mml)を各々溶製した後、表3 の比較例29を除いて、全て510 ℃×8 hrの均質化熱処理 (空気炉) を施した (比較例29は490 ℃×8 hr) 。この均質化熱処理後の鋳塊を熱間鍛造にて、各方向での鍛練比が1.5 以上となるように、150 mm角( 厚み) の角棒および80mm角 (厚み) の角棒としたものを、300mmlの長さに切断してAl合金鍛造材を製造した。このAl合金鍛造材を空気炉で加熱速度200 ℃/hr で昇温し、発明例は表2 、比較例は表3 に各々示す種々の溶体化温度で、かつ時間は共通して6 hrの溶体化処理後、表2 、3 に示す種々の焼入れ温度で水焼入れを行い(400℃から290 ℃の間の平均冷却速度は30000 ℃/ 分以上) 水中で10分保持後に取り出した。
【００４６】
前記厚みが80mmのAl合金鍛造材については、小物部品やピストンなどの残留応力が比較的大きくても良い用途を模擬して、溶体化処理後に30〜45℃の低温の水焼入れ処理し、その後に175 ℃×18hrの人工時効硬化処理を施した、調質T6材とした。
【００４７】
一方、前記厚みが150mm のAl合金鍛造材については、残留応力が問題となる用途を模擬して、溶体化処理後に70〜91℃の温水焼入れして残留応力を低減し、冷間圧縮加工を加えずに、175 ℃×18hrの人工時効硬化処理を施した、調質T61 材とした。また、同じく、残留応力が問題となる用途を模擬して、溶体化処理後に30〜60℃の水焼入れ処理し、表2 、3 に示す冷間圧縮率で冷間圧縮加工を加えて残留応力を低減し、175 ℃×18hrの人工時効硬化処理を施し、調質T652材とした。
【００４８】
これらの調質Al合金鍛造材から供試体を採取し、ASTM837 に規定される小孔穴あけ法にて、残留応力低減の必要がある調質T61 材、調質T652材についてのみ残留応力を測定した。また、供試材の引張特性として、室温での機械的特性 (σB 、耐力、伸び) と、高温特性として、180 ℃×100hr の高温に供試材を暴露した際の、その温度での機械的特性 (σB 、耐力、伸び) 、更に、204 ℃での1000hrクリープ破断強度を測定した。これら試験片は平行部10mmΦ×28mml とした。これらの供試材の引張特性の測定結果を、発明例は表2 、比較例は表3 に各々示す。なお、溶体化処理中にバーニングや、熱間鍛造中に割れが生じたものは、これら特性を測定しなかった。なお、発明例の中から選択して、発明例9 と発明例13および発明例14については、更に、シャルピー衝撃値(J/cm²) と疲労強度(rpm、応力: 190MPa、室温) についても調査した。
【００４９】
以下に、表1 と、表2 、3 から明らかな事項を述べる。
表1 に示すA 〜G およびM 、N の本発明範囲内の化学成分組成を有し、表2 に示す本発明範囲内の鍛造温度、溶体化温度で処理した、発明例1 〜8 までの調質T6材、発明例9 の調質T61 材、発明例10の調質T652材、発明例11、12の調質T6材は、各々室温での耐力が400MPa以上で、室温強度および高温強度、更にクリープ破断強度が高い。なお、発明例の中から選択された、発明例9 、発明例13および発明例14については、シャルピー衝撃値(J/cm²) は各々3.0 、4.5 、4.3 であり、疲労強度(rpm) は各々3.0e6 、5.5e6 、5.8e6 であった。
【００５０】
これに対し、表1 に示すA の本発明範囲内の合金を用いても、表3 に示す鍛造温度が435 ℃と本発明範囲を高めに外れる比較例20 (調質T6材) は、上記発明例1 〜8 までの調質T6材に比して、特にクリープ破断強度が低い。また、表3 に示す鍛造温度が260 ℃と本発明範囲を低めに外れる比較例21 (調質T6材) は、鍛造割れが生じた。したがって、鍛造温度の本発明範囲の意義が裏付けられる。
【００５１】
同じく、表1 に示すA の本発明範囲内の合金を用いても、表3 に示す溶体化処理温度が550 ℃と本発明範囲を高めに外れる比較例22 (調質T6材) と比較例25 (調質T61 材) は溶体化処理中にバーニングが生じた。また、表3 に示す溶体化処理温度が500 ℃と本発明範囲を低めに外れる比較例23 (調質T6材) と比較例26 (調質T61 材) は、上記発明例1 〜9 の調質T6材や調質T61 材同士での比較で、室温強度および高温強度と、特にクリープ破断強度が低い。したがって、溶体化処理温度の本発明範囲の意義が裏付けられる。
【００５２】
また、製造条件が略同じ場合で、合金成分のみが違う場合について、以下に比較する。
発明例の中でも、Agの含有量が比較的少ない合金B を用いた発明例2 は、Agの含有量が下限未満の合金J を用いた比較例17よりも、室温強度および高温強度、更にクリープ破断強度が著しく高い。しかし、よりAgの含有量が多い合金A を用いた発明例1 などよりも、室温強度および高温強度、更にクリープ破断強度が比較的低い。したがって、Agの含有効果と0.05% の下限含有量の意義が裏付けられる。
【００５３】
Si含有量が比較的高い合金C を用いた発明例3 や、Fe含有量が比較的高い合金D を用いた発明例4 、 Cr 含有量が比較的高い合金E を用いた発明例5 、Zr含有量が比較的高い合金F を用いた発明例6 、Mn含有量が比較的高い合金G を用いた発明例7 は、これらの含有量が低い発明例1 に比較して室温強度および高温強度、更にクリープ破断強度が比較的低い。したがって、これら不純物を各々所定量以下に規制する意義が裏付けられる。
【００５４】
Cu含有量が4.1%、5.3%と比較的少ない合金M やN を用いた発明例11、12 (調質T6材) および発明例13、14 (調質T61 材) は、Cu含有量が6.3%と比較的多く、製造条件が各々同じ発明例1 、2(調質T6材) および発明例9(調質T61 材) などに比して、高温強度またはクリープ破断強度などの高温特性が比較的低い。また、Cu含有量が少なすぎる合金H を用いた比較例15 (調質T6材) 、Mg含有量が少なすぎる少なすぎる合金I を用いた比較例16 (調質T6材) は、製造条件が同じ発明例1 、2 などに比較して、特に、高温強度またはクリープ破断強度などの高温特性が著しく低い。Mg含有量が多すぎる合金K を用いた比較例18 (調質T6材) は溶体化処理中にバーニングが生じた。また、比較のために本発明組成範囲を外れたJIS 2618合金 Lを用いた比較例19 (調質T61 材) は発明例に比較して、特に、高温強度またはクリープ破断強度などの高温特性が著しく低い。したがって、これらの結果から本発明成分組成範囲や好ましい成分組成範囲の意義が裏付けられる。
【００５５】
次に、焼入れ温度条件以外の条件が略同じ場合について、焼入れ温度条件の違いについて比較する。調質T6材において焼入れ温度が45℃と高い比較例24、調質T652材において焼入れ温度が60℃と高い比較例28は、調質T6材において焼入れ温度が30℃と低い発明例1 、8 や、調質T652材において焼入れ温度が30℃と低い発明例10の同じ調質材同士で比較して、特に、室温強度や、高温強度、クリープ破断強度などの高温特性が著しく低い。したがって、焼入れ温度が40℃以下と低い方が、室温強度や、高温強度またはクリープ破断強度などの高温特性が高いことが分かる。
【００５６】
また、調質T61 材において、焼入れ温度が70℃と低すぎる比較例27は、他の焼入れ温度が91℃の発明例9 などに比して、残留応力が大きすぎるため、前記した残留応力が問題となる用途では使用できない。
【００５７】
更に、調質T652材において、冷間圧縮率が低すぎる比較例29は、他の調質T652材に比して、室温強度や、高温強度またはクリープ破断強度などの高温特性は高いものの、残留応力が大きすぎるため、前記した残留応力が問題となる用途では使用できない。一方、冷間圧縮率が大きすぎる比較例30は、調質T652材の発明例10に比して、室温強度や、高温強度またはクリープ破断強度などの高温特性は遜色ないものの、残留応力が大きすぎるため、前記した残留応力が問題となる用途では使用できない。したがって、残留応力を除去する冷間圧縮率には前記した適正範囲があることが分かる。また、発明合金A を用いているものの、均質化熱処理温度が490 ℃と本発明範囲よりも低過ぎる比較例31は、均質化不足のために、発明例に比して、室温、高温の引張特性やクリープ破断強度が著しく劣る。
【００５８】
上記各発明例と一部比較例の前記した方法でのミクロ組織観察の結果、発明例のいずれもが、Al合金組織が等軸で、平均結晶粒径が50〜500 μmmの範囲の一定サイズの粒径であり、更に(100) 面上にθ' 相、(110) 面上にΩ相が各々析出していた。これに対して、特に、鍛造温度が上限を越えている前記比較例20では、サブグレインが生じ、等軸再結晶粒も一部存在するものの、再結晶粒が集合体化した粗大な再結晶粒とからなっていた。したがって、室温強度や、高温強度またはクリープ破断強度などの高温特性が同様に低い他の比較例も、前記好ましい本発明のミクロ組織から外れたミクロ組織となっていると考えられる。
【００５９】
【表１】

【００６０】
【表２】

【００６１】
【表３】

【００６２】
【発明の効果】
本発明によれば、冶金的に設計されたAl合金鍛造材を、その設計高温特性通りに再現性良く製造できるAl合金鍛造材の製造方法を提供することができる。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a 2000 series aluminum alloy forging (hereinafter, aluminum is simply referred to as Al), and relates to a method for producing an Al alloy forging excellent in high temperature characteristics (heat resistance and high temperature proof stress).
[0002]
[Prior art]
For aviation and space equipment such as rockets and airplanes, for transportation equipment such as railway vehicles, automobiles and ships, or for machine parts such as engine parts and compressors, specifically for rotating rotors, rotating impellers or pistons, etc. Al alloy forgings with excellent high temperature characteristics are used for Al alloy parts that are used in high temperature environments exceeding 100 ° C. The high temperature characteristics are creep resistance characteristics and high temperature proof stress at the high temperatures.
[0003]
Conventionally, for these so-called heat-resistant Al alloy forgings, AA standard to JIS standard 2000 series (hereinafter simply referred to as 2000 series) Al alloys are used. Examples of this kind of Al alloy include 2219 and 2618. However, these 2000 series Al alloys have a significant decrease in strength when used for a long time at temperatures exceeding 120 ° C.
[0004]
For this reason, in order to improve the creep characteristics and high temperature proof stress in high temperature usage environments exceeding 120 ° C, in recent years, 2519Al alloy (Al-6.1Cu-0.3Mn-0.15Zr-0.1 V) has been developed. In addition, a 2519 (Ag) Al alloy obtained by adding Ag to the 2519Al alloy has been developed. Many Al alloys related to these 2519Al alloys and 2519 (Ag) Al alloys have also been proposed (see, for example, Patent Documents 1 and 2). The present inventors have also proposed a heat-resistant Al alloy material that can guarantee high temperature characteristics with good reproducibility. This content includes Cu: 1.5 to 7.0%, Mg: 0.01 to 2.0%, and optionally, heat resistant Al alloy containing Ag: 0.05 to 0.7%, for θ ′ phase and / or Ω phase, θ ′ The average size of the phase is 120 nm or less and the average interval between the precipitates of the θ ′ phase is 100 nm or less, the average size of the Ω phase is 100 nm or less, and the average interval between the precipitates of the Ω phase is 150 nm or less (See Patent Document 3 and Non-Patent Document 1).
[0005]
[Patent Document 1]
Japanese Patent Laid-Open No. 62-112748 [Patent Document 2]
US Pat. No. 4,610,733 [Patent Document 3]
JP 11-302764 [Non-patent Document 1]
Outline of the 93rd Autumn Meeting of the Japan Institute of Light Metals (October 20, 1997, pp. 233-234)
[0006]
In addition, the application parts that require the high temperature characteristics basically have a thick cylindrical shape or a complicated shape in which a large number of blades are provided around. For this reason, when these parts are manufactured using an Al alloy material, a forged material obtained by hot forging (including cold forging after hot forging) of an aluminum alloy bulk (bulk) ingot is used. Parts are made by cutting. These application parts are required to have high dimensional accuracy and smoothness because they slide or rotate in a narrow space or clearance at high speed. For this reason, in addition to the said high temperature characteristic, the high precision cutting workability, ie, a machinability, is requested | required of Al alloy material used for these uses.
[0007]
For this reason, in order to guarantee the machinability in the cutting work to the high-speed moving parts together with the high temperature characteristics of the heat-resistant Al alloy forgings for high-speed moving parts, the present inventors It has also been proposed that the microstructure of the above has an θ ′ phase and / or an Ω phase and that the crystal grain size be equiaxed recrystallized grains of 500 μm or less (see Patent Document 4).
[0008]
[Patent Document 4]
JP 2000-119786 A [0009]
However, even if Al-alloy forgings with excellent high-temperature properties are metallurgically designed with these technologies, high-temperature artificial age hardening after solution treatment and quenching is performed on the Al-alloy forgings actually produced. Even if it is applied, the yield strength is not improved, the yield strength after artificial age hardening required for this kind of Al alloy forging (heat resistant Al alloy forging) is lowered, and the yield strength at high temperature use may also be lowered. . For this reason, the present inventors pay attention to the influence of the quenching rate after the solution treatment, and the quenching rate (cooling rate) is low (small) with an average cooling rate between 400 ° C and 290 ° C of 30000 ° C / min or less. In particular, it has been proposed that Zr, Cr, and Mn in the forged Al alloy material are restricted to Zr: 0.09% or less, Cr: 0.05% or less, and Mn: 0.6% or less, respectively (Patent Document 5). See).
[0010]
[Patent Document 5]
JP 2001-181771 A [0011]
[Problems to be solved by the invention]
However, in order to be able to manufacture Al alloy forgings that are metallurgically designed to be excellent in high temperature characteristics with high reproducibility according to the design high temperature characteristics, it is still necessary to provide hot forging conditions, solution hardening conditions, etc. Improvement of actual manufacturing conditions was necessary, and development factors remained.
[0012]
The present invention has been made paying attention to such circumstances, and its purpose is to improve manufacturing conditions such as hot forging conditions and solution quenching treatment conditions, and to manufacture with high reproducibility according to the design high temperature characteristics. An object of the present invention is to provide a method for producing a forged Al alloy material.
[0013]
[Means for Solving the Problems]
In order to achieve this object, the gist of claim 1 of the method for producing an aluminum alloy forging according to the present invention includes Cu: 4.0 to 7.0%, Mg: 0.2 to 0.4%, Ag: 0.05 to 0.7%, and the balance aluminum. And a forged aluminum alloy forging material comprising the inevitable impurities, wherein the cast material having this composition is subjected to homogenization heat treatment at a temperature of 500 to 535 ° C., hot forged at a temperature of 280 to 360 ° C. , and then 510 The aluminum alloy forging material has a property that the yield strength at room temperature is 400 MPa or more when it is subjected to solution treatment and quenching treatment at a temperature of ˜545 ° C. and artificial age hardening treatment. In addition,% display of alloy element content means mass%.
[0014]
Here, in order to produce the metallurgically designed Al alloy forging material with high reproducibility according to its design high temperature characteristics, it is doubtful that hot forging conditions and solution hardening conditions are important. Absent. However, the present inventors have found that there are optimum conditions for these hot forging and solution hardening treatments.
[0015]
That is, first, the hot forging temperature needs to be lower than the general hot forging temperature condition range in order to manufacture with good reproducibility according to the design high temperature characteristics. Moreover, in order to produce the solution treatment temperature and the quenching temperature after the solution treatment with good reproducibility according to the design high temperature characteristics, as described later, the treatment should be performed within the above temperature range as found by the present inventors. is required.
[0016]
DETAILED DESCRIPTION OF THE INVENTION
The significance of each requirement of the present invention will be described below.
The manufacturing process itself of the Al alloy forged material in the present invention is basically the same as the conventional one. That is, the ingot is prepared by appropriately selecting an ordinary molten casting method such as a continuous casting rolling method and a semi-continuous casting method (DC casting method) from a molten Al alloy melt adjusted within the component range of the present invention. To manufacture. This ingot is subjected to homogenization heat treatment at a temperature of 500 to 535 ° C. and then hot forged to produce an Al alloy forged material. In addition, as a raw material for forging, you may use the extruded material and rolled material which extruded and rolled the ingot. Here, when the temperature of the homogenization heat treatment is less than 500 ° C., the ingot crystallization product does not dissolve, and the homogenization becomes insufficient. On the other hand, if the temperature of the homogenization heat treatment exceeds 535 ° C., the possibility of burning is increased. Therefore, the temperature of the homogenization heat treatment is in the range of 500 to 535 ° C.
[0017]
Here, the manufacturing conditions and manufacturing means of the Al alloy forged material in the present invention include the hot forging temperature, the solution treatment temperature, the quenching temperature after solution treatment, and the cold compression conditions performed as necessary, as described in detail below. Is basically the same as the conventional method. In other words, it is an advantage of the present invention that the production conditions and production means of the Al alloy forging are not changed greatly.
[0018]
First, the temperature conditions for hot forging are important for producing an Al alloy forged material with good reproducibility according to the design high temperature characteristics. Conventionally, in order to make the microstructure after solution treatment of Al alloy forging material equiaxed crystal grains, even if well-known forging means such as free forging and die forging (forging forging) are used alone or in combination, as appropriate In addition, the hot forging temperature was set to about 380 to 430 ° C. This is because when the hot forging temperature is low, the structure of the Al alloy forged material is likely to be locally mixed grains and the high temperature characteristics are deteriorated.
[0019]
In this regard, in the present invention, the hot forging temperature is set to a temperature range of 280 to 360 ° C. below the recrystallization temperature. When the hot forging temperature exceeds 360 ° C., coarse grains are likely to occur in the Al alloy forging material within the component range of the present invention. For this reason, the high temperature characteristic of Al alloy forging material falls, and the Al alloy forging material excellent in the high temperature characteristic cannot be manufactured with good reproducibility. On the other hand, if the hot forging temperature is less than 280 ° C., cracking is likely to occur during hot forging, and forging itself becomes difficult.
[0020]
In the present invention, even if the temperature of hot forging is set to 280 to 360 ° C., the Al alloy forging material within the component range of the present invention can be obtained after the tempering of the Al alloy forging material by appropriate solution treatment and quenching treatment. The structure becomes equiaxed grains, not mixed grains.
[0021]
The microstructure of the Al alloy forging material is also affected by the forging ratio in hot forging. Therefore, in the case of an Al alloy forged material, in order to make the microstructure into equiaxed grains, it is preferable that the appropriate hot forging forging ratio is 1.5 or more. If the forging ratio is less than 1.5, the structure of the Al alloy forged material tends to be mixed grains. Further, it is more preferable that the training direction is not limited to one direction but at least two different directions, and the training ratio in each direction is 1.5 or more.
[0022]
Next, solution treatment and quenching treatment will be described. In this solution treatment and quenching treatment, JIS-H-4140, AMS-H-6088, etc. specified in order to re-dissolve soluble intermetallic compounds and suppress reprecipitation during cooling as much as possible. It is preferable to carry out within the conditions. However, even if heat treatment is performed according to standards such as AMS-H-6088, if the solution treatment temperature is too high, burning occurs and the mechanical properties are significantly reduced. When the solution treatment temperature is below the lower limit, the yield strength at room temperature after the artificial age hardening treatment does not exceed 400 MPa, and the solution treatment itself becomes difficult. Therefore, the upper limit of the solution treatment temperature is 545 ° C., and the lower limit is 510 ° C.
[0023]
Here, in applications such as small parts up to about φ100 mm and pistons, even if the residual stress is relatively large, products that do not cause any processing problems such as cutting are subjected to artificial age hardening after solution treatment and quenching. It is desirable to apply tempered T6. In this case, it is desirable that the quenching temperature is 40 ° C. or lower in order to obtain high strength characteristics and high temperature characteristics even if the residual stress becomes relatively large. Moreover, when this quenching temperature is high, it becomes difficult to make the yield strength at room temperature after the artificial age hardening treatment be 400 MPa or more.
[0024]
On the other hand, in a large product such as a rotor, a high residual stress exceeding 10 kgf / mm ² is generated on the product surface because the cooling rate of the product surface and the central part is greatly different during quenching. When such a high residual stress is generated, a large distortion occurs during the cutting of the product, and precise cutting becomes extremely difficult. In the worst case, breakage such as cracking due to residual stress may occur during cutting. For example, even if no breakage such as cracking occurs during cutting, the product starts from an intermetallic compound such as a crystallized substance remaining in the material or from a slight surface flaw generated during product transportation. During long-term use, cracks tend to propagate and grow, which may lead to final fracture. Therefore, for products where residual stress is a problem, such as rotors, the water quenching temperature after solution treatment should be a relatively high temperature of 90 ° C. or higher in order to remove or reduce the residual stress to preferably 3.0 kgf / mm ² or less. Thereafter, an artificial age hardening treatment is preferably performed to obtain a tempered T61 material.
[0025]
In addition, depending on the application, there is a product whose residual stress is strictly controlled regardless of the size of the product. For such products, in order to minimize the residual stress, cold compression or cold working is applied to remove or reduce the residual stress to preferably 3 kgf / mm ² or less, and an artificial age hardening treatment is applied. Preferably, the material is T652. In these products, the residual stress is preferably removed or reduced to 3 kgf / mm ² or less, and in order to obtain high strength characteristics and high temperature characteristics, the quenching temperature is preferably 40 ° C. or less. When the quenching temperature is high, it becomes difficult to make the yield strength at room temperature after the artificial age hardening treatment be 400 MPa or more. If the amount of cold compression (cold working) of cold working or cold working is small, a sufficient residual stress reducing effect cannot be obtained. On the other hand, when the amount of cold compression is large, the amount of precipitation of the θ ′ phase increases during the artificial age hardening treatment or during use at a high temperature, so that the proof stress tends to decrease. Therefore, it is preferable that the cold compression (processing) has a compression (processing) rate of 1 to 5%.
[0026]
Thereafter, these Al alloy forgings are processed into the above-mentioned application parts. Of course, after the Al alloy forged material is processed into the above-mentioned application product, solution treatment, quenching treatment, cold compression, artificial age hardening treatment, and the like may be appropriately performed.
[0027]
As a furnace used for tempering (heat treatment) such as solution treatment and quenching, a batch furnace, a continuous annealing furnace, a molten salt bath furnace, an oil furnace, or the like can be used as appropriate. Further, as a cooling means at the time of quenching, means such as Yukon quell chant, water immersion, hot water immersion, boiling water immersion, water injection, and air injection can be appropriately selected.
[0028]
The microstructure of the Al alloy forging of the present invention thus obtained has a fine structure of almost constant size of 500 μm or less, preferably in the range of 10 to 500 μm, more preferably in the range of 50 to 300 μm. Recrystallized grains (equal axis recrystallized grains). And as seen in the mixed grain structure, a group of fine recrystallized grains (or sub-crystal grains) having a grain size of 1 μm or less, a coarse recrystallized grain of several mm to several cm, Alternatively, there is no remaining ingot structure, and it has both high temperature characteristics such as good creep characteristics and machinability.
[0029]
However, the preferred structure of the equiaxed recrystallized grains in the present invention does not necessarily mean a structure having only 100% of the equiaxed recrystallized grains of a certain size, and the high temperature such as the machinability and creep rupture strength. Mixing of a cast structure or a mixed grain structure within a range that does not deteriorate the characteristics is allowed. For example, fine recrystallized grains (or sub-crystal grains) with a grain size of 1 μm or less have high temperature characteristics such as machinability and creep rupture strength, even if single crystal grains are dispersed individually. Does not decrease. However, when these are assembled or assembled in a form of sticking to each other, the machinability and the high temperature characteristics are lowered. Therefore, from this point, the area ratio of aggregated fine recrystallized grains of 1 μm or less in the microstructure after solution treatment is preferably 10% or less.
[0030]
Incidentally, the specification of equiaxed recrystallized grains and the presence or absence of mixed grain structure as used in the present invention can be observed or measured with an optical microscope of 50 to 400 times by micro-etching a sample by electrolytic etching or the like.
[0031]
Next, in the Al alloy forging structure of the present invention, in order to further enhance the high temperature characteristics such as high temperature proof stress and creep rupture strength, the range of 160 to 190 ° C. × 7 to 60 hours after solution treatment and quenching treatment It is preferable to precipitate the θ ′ phase precipitated on the (100) plane of the Al alloy and the Ω phase precipitated on the (111) plane. Without these precipitations due to the artificial age hardening treatment, even when the artificial age hardening treatment is performed, the high-temperature proof stress at a temperature such as 180 ° C. becomes low.
[0032]
The precipitation state of the θ ′ phase and the Ω phase in the Al alloy forging material structure can be identified by observing the structure with a transmission electron microscope (TEM) of 50000 times.
[0033]
The chemical component composition in the Al alloy forged material of the present invention will be described. The chemical composition of the Al alloy of the present invention is basically good as a component standard for Al alloys such as 2519 or 2618 and 2519 (Ag) -based Al alloys obtained by adding Ag to 2519, but more specific applications and required characteristics Depending on the above, it can be appropriately selected from the component composition ranges described below. First, active elements will be described.
[0034]
(Cu: 4.0-7.0%)
Cu is a basic component of the Al alloy forging material of the present invention. By both the solid solution strengthening and precipitation strengthening effects, it has the normal-temperature and high-temperature creep characteristics and high-temperature proof stress that are mainly required in the present invention application of the Al alloy forging material. It is essential to secure. More specifically, Cu, as described above, causes the θ ′ phase and the Ω phase to be finely and densely deposited on the (100) surface and (111) surface of the Al alloy during the high-temperature artificial age hardening treatment. , Improve the strength of Al alloy forgings after artificial age hardening. This effect is exhibited at 4.0% or more. When the Cu content is less than 4.0%, the above-described effect is small, and sufficient creep characteristics and high-temperature proof stress of the Al alloy forging material at normal temperature and high temperature cannot be obtained. On the other hand, if the Cu content exceeds 7.0%, the strength becomes too high, and the forgeability of the Al alloy forged material decreases. Therefore, the Cu content is in the range of 4.0 to 7.0%.
[0035]
(Mg: 0.2-0.4%)
Mg, as well as Cu, is essential to ensure sufficient creep properties and high-temperature proof strength of Al alloy forgings at normal and high temperatures mainly by the effects of both solid solution strengthening and precipitation strengthening. More specifically, Mg, like Cu, has a fine and high density of the θ 'phase and Ω phase on the (100) and (111) surfaces of the Al alloy forging during high temperature artificial age hardening. Precipitate and improve the strength of the Al alloy forging material after artificial age hardening treatment. This effect is exhibited at 0.2% or more. If the Mg content is less than 0.21%, this effect is not exhibited, and sufficient creep characteristics and high temperature proof stress at room temperature and high temperature of the Al alloy forging cannot be obtained. On the other hand, if the Mg content exceeds 0.4%, the strength becomes too high, and there is a high possibility that cracks referred to as burning will occur during solution treatment or that forgeability will be reduced. Therefore, the Mg content is in the range of 0.2 to 0.4%.
[0036]
(Ag: 0.05-0.7%)
Ag forms a fine and uniform Ω phase in the Al alloy forging, and the width of the area where no precipitate phase exists (PFZ: solute-depleted precipitate free zone) is extremely narrow. Indispensable for improving normal temperature and high temperature strength. If the Ag content is less than 0.05%, this effect is not obtained. On the other hand, if the Ag content exceeds 0.7%, the effect is saturated. Therefore, the Ag content is in the range of 0.05 to 0.7%.
[0037]
(V: 0.15% or less)
V, like Zr, Cr, and Mn, precipitates Al-V-based dispersed particles, which are thermally stable compounds in the Al alloy forging structure during homogenization heat treatment, and these dispersed particles are recrystallized after recrystallization. This has the effect of preventing the grain boundary from moving, thus preventing the grain coarsening. As a result, the microstructure of the Al alloy forging is made into a fiber structure, and there is an effect of improving the normal temperature strength and the high temperature strength. The action of coarsely depositing the stable phase is relatively small compared to Zr, Cr, and Mn. Therefore, in order to make the crystal grain size finer to 500 μm or less for the purpose of more reliably ensuring the machinability and high temperature characteristics of the Al alloy forging, in a preferred embodiment of the present invention, V is 0.15% or less. Is preferably contained selectively. If the V content is less than 0.05%, these effects are difficult to obtain. On the other hand, if the V content exceeds 0.15%, a coarse insoluble intermetallic compound is likely to be formed during melting and casting, which causes molding defects and breakage. Therefore, V is 0.15% or less, preferably 0.05% to 0.15%.
[0038]
The elements that are preferably regulated will be described below.
Zr, Cr, and Mn are dispersed in the Al-Zr, Al-Cr, and Al-Mn systems, which are thermally stable compounds in the Al alloy forging material structure, respectively, during the homogenization heat treatment. Precipitate particles. The dispersed particles have an effect of improving the normal temperature strength and the high temperature strength by forming the microstructure of the Al alloy forged material into a fiber structure.
[0039]
However, when the average cooling rate between 400 ° C and 290 ° C is slowed to 30000 ° C / min or less in the quenching process after solution treatment, if these Zr, Cr and Mn are contained, the solution treatment In the subsequent quenching process, a stable phase such as AlCu ₂ is coarsely deposited around the Al—Cr, Al—Zr, and Al—Mn dispersed particles in the quenching process. As a result, even if the artificial age-hardening treatment at a high temperature is performed next, the yield strength at a high temperature such as 310 MPa or more cannot be obtained after being used at a temperature of 120 ° C. for 100 hours. Therefore, in order to reduce the quenching sensitivity of the Al alloy forged material, it is preferable that Zr: 0.09% or less, Cr: 0.05% or less, and Mn: 0.8% or less are preferably controlled.
[0040]
Fe is preferably regulated to 0.15% or less. However, there is also contamination from scraps and the like, and there is an effect of improving the high temperature characteristics of the Al alloy forging material, so a content of 0.15% is allowed. If the content exceeds 0.15%, an insoluble intermetallic compound is formed, which tends to cause molding defects and breakage.
[0041]
Si combines with Mg to form Mg ₂ Si and Al-Fe-Si based crystals in the Al alloy forging structure. For this reason, Mg is necessary to precipitate the θ 'phase and Ω phase during the high-temperature artificial age hardening treatment and improve the strength of the Al alloy forging after the artificial age hardening treatment. The strength of the Al alloy material after the treatment is lowered. Since the Mg content is originally lower than that of Cu, the influence of Si is large. In addition, most of the crystallized product is dissolved by the solution treatment, but if excessive Mg ₂ Si is formed, it remains in the solution treatment and becomes the starting point of fracture, so that the formability is lowered. . Therefore, Si is preferably regulated to 0.1% or less.
[0042]
In addition, Ti refines crystal grains, but if added excessively, a coarse intermetallic compound is formed and becomes a starting point of fracture at the time of molding, so that formability is lowered. Therefore, Ti is allowed to be contained up to 0.1% or less.
[0043]
Therefore, in a preferred embodiment of the present invention, the following elements in the Al alloy forging alloy are used in order to prevent the yield strength after artificial age hardening of the Al alloy forging material from being lowered and the yield strength at the time of high temperature use from being lowered. Are preferably regulated to Si: 0.1% or less, Fe: 0.15% or less, Zr: 0.09% or less, Cr: 0.05% or less, Mn: 0.8% or less, and Ti: 0.1% or less.
[0044]
For elements other than the above, such as Zn, Ni, B, etc., inclusion in a range that does not impair the high temperature characteristics and other characteristics of the Al alloy forging according to the present invention or the upper limit of about 2000 series Al alloy is allowed. The
[0045]
【Example】
Next, examples of the present invention will be described. After melting the aluminum alloy ingots (500 mm φ × 2000 mml) of chemical composition within the scope of the present invention of A to G and chemical composition outside the scope of the present invention of H to L shown in Table 1, respectively, Except for Comparative Example 29, all were subjected to a homogenization heat treatment (air furnace) of 510 ° C. × 8 hr (Comparative Example 29 was 490 ° C. × 8 hr). The ingot after this homogenization heat treatment is hot forged into 150 mm square (thickness) square bars and 80 mm square (thickness) square bars so that the forging ratio in each direction is 1.5 or more. Was cut into a length of 300 mml to produce an Al alloy forged material. The aluminum alloy forging was heated in an air furnace at a heating rate of 200 ° C./hr.Invention examples are shown in Table 2 and comparative examples are shown in Table 3. After solution treatment, water quenching was performed at various quenching temperatures shown in Tables 2 and 3 (the average cooling rate between 400 ° C. and 290 ° C. was 30000 ° C./min or more), and the sample was taken out after being kept in water for 10 minutes.
[0046]
For the Al alloy forged material having a thickness of 80 mm, a low-temperature water quenching process of 30 to 45 ° C. is performed after the solution treatment, simulating the use where residual stresses such as small parts and pistons may be relatively large, and thereafter The tempered T6 material was subjected to an artificial age hardening treatment at 175 ° C. for 18 hours.
[0047]
On the other hand, for Al alloy forgings with a thickness of 150mm, the residual stress is reduced by hot water quenching at 70 to 91 ° C after solution treatment, simulating applications where residual stress is a problem. Without any addition, the tempered T61 material was subjected to artificial age hardening at 175 ° C x 18hr. Similarly, the residual stress is simulated by water quenching at 30 to 60 ° C after solution treatment, and cold compression is applied at the cold compressibility shown in Tables 2 and 3 simulating applications where residual stress is a problem. And subjected to artificial age hardening at 175 ° C x 18 hr to obtain a tempered T652 material.
[0048]
Specimens were collected from these tempered Al alloy forgings, and the residual stress was measured only for tempered T61 and tempered T652 materials that required residual stress reduction by the small hole drilling method specified in ASTM 837. . In addition, mechanical properties at room temperature (σB, proof stress, elongation) as tensile properties of the test material and mechanical properties at that temperature when the test material was exposed to a high temperature of 180 ° C x 100 hr as high temperature properties. Characteristics (σB, proof stress, elongation) and 1000 hr creep rupture strength at 204 ° C. were measured. These test pieces had a parallel portion of 10 mmΦ × 28 mml. The measurement results of the tensile properties of these test materials are shown in Table 2 for the inventive examples and Table 3 for the comparative examples. In addition, those in which cracking occurred during burning or hot forging during solution treatment were not measured. In addition, selected from the inventive examples, for the inventive examples 9, 13 and 14, the Charpy impact value (J / cm ² ) and fatigue strength (rpm, stress: 190 MPa, room temperature) investigated.
[0049]
The matters that are clear from Table 1 and Tables 2 and 3 are described below.
Inventive Examples 1 to 8 having chemical composition compositions A to G and M and N shown in Table 1 within the scope of the present invention and treated at forging temperature and solution temperature within the scope of the present invention shown in Table 2. The tempered T6 material, the tempered T61 material of Invention Example 9, the tempered T652 material of Invention Example 10, and the tempered T6 material of Invention Examples 11 and 12, each had a proof stress at room temperature of 400 MPa or more, room temperature strength and high temperature strength Furthermore, the creep rupture strength is high. Inventive Example 9, Inventive Example 13 and Inventive Example 14, which were selected from the inventive examples, had Charpy impact values (J / cm ² ) of 3.0, 4.5 and 4.3, respectively, and fatigue strength (rpm) was They were 3.0e6, 5.5e6, and 5.8e6, respectively.
[0050]
In contrast, even when an alloy within the scope of the present invention of A shown in Table 1 is used, the forging temperature shown in Table 3 is 435 ° C., which is a comparative example 20 (tempered T6 material) that deviates from the scope of the present invention. Compared with the tempered T6 material of Invention Examples 1 to 8, the creep rupture strength is particularly low. Further, forging cracks occurred in Comparative Example 21 (tempered T6 material), which has a forging temperature of 260 ° C. shown in Table 3 and deviates from the scope of the present invention. Therefore, the significance of the forging temperature in the present invention is supported.
[0051]
Similarly, even when an alloy within the scope of the present invention of A shown in Table 1 is used, the solution treatment temperature shown in Table 3 is 550 ° C., which is far from the scope of the present invention, and Comparative Example 22 (tempered T6 material) and Comparative Example No. 25 (tempered T61 material) burned during solution treatment. In addition, Comparative Example 23 (tempered T6 material) and Comparative Example 26 (tempered T61 material), which have a solution treatment temperature of 500 ° C. shown in Table 3 and deviate from the scope of the present invention, are prepared according to the above-described inventive examples 1 to 9. Compared with high quality T6 material and tempered T61 material, room temperature strength and high temperature strength, especially creep rupture strength are low. Therefore, the significance of the scope of the present invention of the solution treatment temperature is supported.
[0052]
Further, the case where the manufacturing conditions are substantially the same and only the alloy components are different will be compared below.
Among Invention Examples, Invention Example 2 using an alloy B having a relatively low Ag content is higher in room temperature strength and high temperature strength and further creeping than Comparative Example 17 using an alloy J having an Ag content less than the lower limit. Breaking strength is extremely high. However, the room temperature strength, high temperature strength, and creep rupture strength are relatively lower than those of Invention Example 1 using the alloy A having a higher Ag content. Therefore, the significance of the Ag content effect and the lower limit content of 0.05% is supported.
[0053]
Invention Example 3 using an alloy C having a relatively high Si content, Invention Example 4 using an alloy D having a relatively high Fe content, Invention Example 5 using an alloy E having a relatively high Cr content, Zr Invention Example 6 using an alloy F having a relatively high content, and Invention Example 7 using an alloy G having a relatively high Mn content, are compared with Invention Example 1 having a low content of these materials at room temperature strength and high temperature strength. Furthermore, the creep rupture strength is relatively low. Therefore, the significance of regulating these impurities to a predetermined amount or less is supported.
[0054]
Invention Examples 11 and 12 (tempered T6 material) and Invention Examples 13 and 14 (tempered T61 material) using alloys M and N having a relatively low Cu content of 4.1% and 5.3% have a Cu content of 6.3. Compared to Invention Examples 1 and 2 (tempered T6 material) and Invention Example 9 (tempered T61 material), which have relatively many manufacturing conditions, the high temperature properties such as high temperature strength or creep rupture strength are compared. Low. Moreover, Comparative Example 15 (tempered T6 material) using an alloy H with too little Cu content, and Comparative Example 16 (tempered T6 material) using an alloy I with too little Mg content have manufacturing conditions Compared to the same Invention Examples 1 and 2, etc., the high temperature characteristics such as high temperature strength or creep rupture strength are particularly low. In Comparative Example 18 (tempered T6 material) using an alloy K having too much Mg content, burning occurred during the solution treatment. Further, for comparison, Comparative Example 19 (tempered T61 material) using JIS 2618 Alloy L outside the composition range of the present invention has particularly high temperature characteristics such as high temperature strength or creep rupture strength as compared with the inventive examples. Remarkably low. Therefore, these results support the significance of the component composition range of the present invention and the preferred component composition range.
[0055]
Next, when the conditions other than the quenching temperature condition are substantially the same, the difference in the quenching temperature condition is compared. Comparative Example 24, which has a high quenching temperature of 45 ° C. in the tempered T6 material, and Comparative Example 28, which has a high quenching temperature of 60 ° C. in the tempered T652, has a low quenching temperature of 30 ° C. in the tempered T6 material. Compared with the same tempered material of Invention Example 10 where the quenching temperature of the tempered T652 material is as low as 30 ° C., the room temperature strength, the high temperature strength, the creep rupture strength, etc. are particularly low. Therefore, it can be seen that the lower the quenching temperature is 40 ° C. or lower, the higher the high temperature characteristics such as room temperature strength, high temperature strength or creep rupture strength.
[0056]
Further, in the tempered T61 material, the comparative stress 27 in which the quenching temperature is too low as 70 ° C. is too large compared with the invention example 9 in which the other quenching temperature is 91 ° C. Cannot be used in problematic applications.
[0057]
Further, in the tempered T652 material, the comparative example 29 in which the cold compressibility is too low is high in room temperature strength, high temperature strength or creep rupture strength, etc. Since the stress is too large, it cannot be used in applications where the residual stress is a problem. On the other hand, Comparative Example 30 with a too high cold compressibility is comparable to Invention Example 10 of the tempered T652 material, although the high temperature characteristics such as room temperature strength, high temperature strength or creep rupture strength are comparable, but the residual stress is large. Therefore, it cannot be used in applications where the above-described residual stress is a problem. Therefore, it can be seen that there is an appropriate range as described above for the cold compressibility for removing the residual stress. Further, although the alloy A is used, the homogenization heat treatment temperature is 490 ° C., which is too lower than the scope of the present invention, in Comparative Example 31, because of insufficient homogenization. Properties and creep rupture strength are extremely inferior.
[0058]
As a result of microstructural observation by the above-described methods of each of the above inventive examples and some comparative examples, all of the inventive examples have a constant size in which the Al alloy structure is equiaxed and the average crystal grain size is in the range of 50 to 500 μmm. Further, the θ ′ phase was precipitated on the (100) plane and the Ω phase was precipitated on the (110) plane. On the other hand, in the comparative example 20 in which the forging temperature exceeds the upper limit, subgrains occur and some of the equiaxed recrystallized grains exist, but the coarse recrystallized grains are aggregated. It consisted of grains. Therefore, it is considered that other comparative examples having low temperature properties such as room temperature strength, high temperature strength, or creep rupture strength are also microstructures that deviate from the preferred microstructure of the present invention.
[0059]
[Table 1]

[0060]
[Table 2]

[0061]
[Table 3]

[0062]
【The invention's effect】
ADVANTAGE OF THE INVENTION According to this invention, the manufacturing method of the Al alloy forging material which can manufacture the Al alloy forging material designed metallurgically according to the design high temperature characteristic with sufficient reproducibility can be provided.

Claims

Cu: 4.0-7.0%, Mg: 0.2-0.4%, Ag: 0.05-0.7%, a method for producing an aluminum alloy forging comprising the balance aluminum and unavoidable impurities, comprising 500 parts of this composition. Aluminum alloy forging after homogenization heat treatment at a temperature of 〜 535 ℃, hot forging at a temperature of 280 〜 360 ℃ , followed by solution treatment and quenching at a temperature of 510 〜 545 ℃ and artificial age hardening treatment A method for producing an aluminum alloy forging material excellent in high temperature characteristics, characterized in that the material has a property that the yield strength at room temperature is 400 MPa or more.

The method for producing an aluminum alloy forged material excellent in high temperature characteristics according to claim 1, wherein the aluminum alloy forged material further contains V: 0.15% or less.

The aluminum alloy forging material includes the following elements in the alloy: Si: 0.1% or less, Fe: 0.15% or less, Zr: 0.09% or less, Cr: 0.05% or less, Mn: 0.8% or less, Ti: 0.1% or less The method for producing an aluminum alloy forging material excellent in high-temperature characteristics according to claim 1 or 2, wherein the forging material is excellent in high temperature characteristics.

The aluminum alloy forging material according to any one of claims 1 to 3, wherein the aluminum alloy forging material is a tempered T6 material having a quenching temperature of 40 ° C or less after the solution treatment . Manufacturing method .

The aluminum alloy forging material according to any one of claims 1 to 3, wherein the aluminum alloy forging material is a tempered T61 material having a quenching temperature of 90 ° C or higher after the solution treatment . Manufacturing method .

The aluminum alloy forged material is a tempered T652 material that has been cold-compressed at a cold compressibility of 1 to 5% after a quenching treatment in which the quenching temperature after the solution treatment is 40 ° C or lower. 4. A method for producing an aluminum alloy forging material having excellent high temperature characteristics according to any one of 3 above.