JP2004084058A - Method for producing aluminum alloy forging for transport structural material and aluminum alloy forging - Google Patents

Method for producing aluminum alloy forging for transport structural material and aluminum alloy forging Download PDF

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Publication number
JP2004084058A
JP2004084058A JP2003046059A JP2003046059A JP2004084058A JP 2004084058 A JP2004084058 A JP 2004084058A JP 2003046059 A JP2003046059 A JP 2003046059A JP 2003046059 A JP2003046059 A JP 2003046059A JP 2004084058 A JP2004084058 A JP 2004084058A
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Japan
Prior art keywords
forging
flash
aluminum alloy
forged
crystal grain
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JP2003046059A
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Japanese (ja)
Inventor
Yoshiya Inagaki
稲垣 佳也
Yasuaki Watanabe
渡辺 泰彰
Manabu Nakai
中井 学
Norifumi Hosoda
細田 典史
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Kobe Steel Ltd
株式会社神戸製鋼所
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Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for producing an aluminum alloy forging for a transport structural material in which, even if hot forging ratios are remarkably different between a product part and a flash part, the fining of crystal grains is possible in both the parts, and to provide an aluminum alloy forging. <P>SOLUTION: An Al-Mg-Si based aluminum alloy casting comprising 0.6 to 1.8% Mg and 0.4 to 1.8% Si, and further comprising one or more kinds of metals selected from 0.01 to 0.9% Mn, 0.01 to 0.25% Cr and 0.01 to 0.20% Zr is subjected to homogenizing heat treatment, and is thereafter subjected to hot forging at a hot forging starting temperature of 450 to 570°C and a final hot forging finishing temperature of ≥360°C. In the crystal grain size in a vertical direction to the parted face 4 of the product part 2 and the flash part 5a in the forging 1 after solution and quenching treatment and artificial age hardening treatment, the average crystal grain size in the product part 2 is controlled to ≤300 μm, and the maximum crystal grain size in the flash part 5a is controlled to ≤400 μm, respectively. <P>COPYRIGHT: (C)2004,JPO

Description

【0001】
【発明の属する技術分野】
本発明は、高強度、高靱性であって、耐応力腐食割れ性などの耐食性にも優れた輸送機構造材用アルミニウム合金鍛造材の製造方法 (以下、アルミニウムを単にAlとも言う) およびアルミニウム合金鍛造材に関するものである。
【0002】
【従来の技術】
周知の通り、車両、船舶、航空機、自動二輪あるいは自動車などの輸送機の構造材乃至部品用、特にアッパーアーム、ロアーアームなどの足回り部品として、AA乃至JIS 6000系(Al−Mg−Si 系) などのAl合金鍛造材が使用されている。6000系Al合金鍛造材は、高強度で高靱性で耐食性にも比較的優れている。また、6000系Al合金自体も、合金元素量が少なく、スクラップを再び6000系Al合金溶解原料として再利用しやすい点で、リサイクル性にも優れている。
【0003】
これら6000系Al合金鍛造材は、Al合金鋳造材を均質化熱処理後、メカニカル鍛造、油圧鍛造などの熱間鍛造(型鍛造)を行い、その後、溶体化および焼き入れ処理と人工時効硬化処理を行う、調質処理が施されて製造される。なお、鍛造素材には、鋳造材を一旦押出した押出材が用いられることもある。
【0004】
近年、これら輸送機の構造材においても、より薄肉化させた上での高強度化や高靱性化が求められている。このため、Al合金鋳造材やAl合金鍛造材のミクロ組織を改善することが種々行われている。
【0005】
例えば、6000系Al合金鋳造材の晶析出物 (晶出物や析出物) の平均粒径を8 μm 以下と小さくし、かつデンドライト二次アーム間隔(DAS) を40μm 以下と細かくして、Al合金鍛造材をより高強度で高靱性化することが提案されている (例えば特許文献1、2参照) 。
また、6000系Al合金鍛造材の結晶粒内や粒界の晶出物や晶析出物の平均粒径や平均間隔などを制御することで、Al合金鍛造材をより高強度で高靱性化することが提案されている (例えば特許文献3、4参照) 。これらの制御は、粒界腐食や応力腐食割れなどに対しても高耐食性化できる。
【0006】
【特許文献1】
特開平07−145440 号公報
【特許文献2】
特開平06−256880 号公報
【特許文献3】
特開2000−144296 号公報
【特許文献4】
特開2001−107168 号公報
【0007】
しかし、これら6000系Al合金鍛造材には、上記鍛造および溶体化処理工程において、加工組織が再結晶して粗大結晶粒が発生する傾向がある。これら粗大結晶粒が発生した場合、上記ミクロ組織を制御しても、高強度化や高靱性化が果たせず、また、耐食性も低下する。
この粗大結晶粒の発生を抑制するため、従来から、Mn、Zr、Crなどの結晶粒微細化効果を有する遷移元素を添加した上で、450 〜570 ℃の比較的高温の温度で熱間鍛造を開始することが知られている (例えば特許文献5参照) 。
【0008】
【特許文献5】
特開平5−247574号公報
【0009】
【発明が解決しようとする課題】
しかし、前記鍛造開始温度を450 〜570 ℃の比較的高温としても、複数回の鍛造工程が再加熱無しあるいは再加熱有りなどで行われる熱間鍛造では、鍛造終了時の鍛造材の温度が比較的低温となることも大いにあり得る。そして、鍛造終了時の鍛造材の温度が比較的低温となった場合、特に、後述するフラッシュ部において、加工組織が再結晶した粗大結晶粒が発生する可能性がある。
【0010】
また、前記足回り部品などの型鍛造品では、製品部とフラッシュ部とが必然的に存在し、このような製品部とフラッシュ部とでは、熱間鍛造時の加工率が大きく異なる。例えば、通常、再加熱無しで複数回行われる熱間鍛造において、前記製品部の加工率は複数回の鍛造の合計でも80% 以下、前記フラッシュ部の加工率は複数回の鍛造の合計で80% 以上と異なる。このような製品部とフラッシュ部との加工率が異なる熱間鍛造では、前記熱間鍛造開始温度を450 〜570 ℃の比較的高温としても、加工歪みをより加えられたフラッシュ部では、溶体化処理工程において、特に、加工組織が再結晶して粗大結晶粒が発生しやすい。
【0011】
これに対し、フラッシュ部の加工歪みを低減するため、フラッシュ部の加工率を下げようとすると、必然的に製品部の方の加工率も下がり、鋳造組織が残留して、強度や靱性が低下したり、形状精度が出なくなる可能性も出てくる。
【0012】
この問題は、本発明者らが先に提案した特願2002−76567号においても生じうる。即ち、この発明では、Mg:0.6〜1.8%、Si:0.4〜1.8%を含み、更に、Mn:0.01 〜0.6%、Cr:0.1〜0.2%およびZr:0.1〜0.2%の一種または二種以上を含み、残部Al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であって、人工時効処理後のアルミニウム合金鍛造材の、0.2%耐力が300MPa以上およびシャルピー衝撃値が10J/cm2 以上であり、更に、製品とフラッシュとの切断面の組織における、型割り面に垂直な方向の結晶粒径であって、この結晶粒径の内の最大のものが400 μm 以下とすることを特徴としている。
【0013】
しかし、この発明でも、好ましい熱間鍛造開始温度は、実施例ともに、350 〜450℃の比較的低い温度であり、複数回の鍛造工程が再加熱無しで行われる熱間鍛造の場合には、鍛造終了時の鍛造材の温度が比較的低温となることも大いにあり得る。この結果、前記フラッシュ部、更には製品部において、加工組織が再結晶した粗大結晶粒が発生する可能性をなお有している。
【0014】
この様な事情に鑑み、本発明の目的は、製品部とフラッシュ部とで熱間鍛造加工率が大きく異なっても、製品部とフラッシュ部ともに、結晶粒の微細化が可能な輸送機構造材用アルミニウム合金鍛造材の製造方法およびアルミニウム合金鍛造材を提供しようとするものである。
【0015】
【課題を解決するための手段】
この目的を達成するために、本発明輸送機構造材用アルミニウム合金鍛造材の製造方法の要旨は、製品部とフラッシュ部とを有するアルミニウム合金鍛造材の製造方法であって、Mg:0.6〜1.8%、Si:0.4〜1.8%を含み、更に、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上を含むAl−Mg−Si系アルミニウム合金鋳造材を、均質化熱処理後に熱間鍛造を行うに際し、前記製品部の熱間鍛造開始温度を450 〜570 ℃とするとともに、製品部の最終熱間鍛造終了温度を360 ℃以上とし、溶体化および焼き入れ処理と人工時効硬化処理後の鍛造材の、前記製品部とフラッシュ部との型割り面に対し垂直方向の結晶粒径の内、前記製品部では平均結晶粒径を300 μm 以下とするとともに、前記フラッシュ部では最大の結晶粒径を400 μm 以下と各々することである。
【0016】
本発明では、前記製品部の加工率が、複数回の鍛造の合計で80% 以下、前記フラッシュ部の加工率が複数回の鍛造の合計で80% 以上となるような製品部とフラッシュ部との加工率が異なる熱間鍛造に際し、製品部の熱間鍛造開始温度を450 〜570 ℃とするとともに、製品部の最終回次の熱間鍛造の終了温度を360 ℃以上とする。そして、製品部の最終回次の熱間鍛造の開始温度と終了温度とをこのように規定することで、製品部の結晶粒径の粗大化防止とともに、前記フラッシュ部の方の結晶粒径の粗大化をも防止乃至抑制できる。なお、本発明で、熱間鍛造の終了温度や開始温度、あるいは、均質化熱処理、鍛造後の調質処理等で規定する温度は、全て鋳造材や鍛造材製品部の外表面の温度である。
例えば、製品部の熱間鍛造の開始温度を450 〜570 ℃としても、製品部の熱間鍛造の終了温度が360 ℃よりも低くなった場合には、Mn、Zr、Crなどの結晶粒微細化効果を有する遷移元素を添加したとしても、調質後の鍛造材の、前記製品部とフラッシュ部との型割り面に対し垂直方向の結晶粒径の内、前記製品部では平均結晶粒径を300 μm 以下とするとともに、前記フラッシュ部では最大の結晶粒径を400 μm 以下と各々することができない。この結果、アルミニウム合金鍛造材を高強度化、高靱性化できなくなり、特に、フラッシュ部の高耐食性化ができない。
【0017】
また、前記目的を達成するための、本発明輸送機構造材用アルミニウム合金鍛造材の要旨は、製品部とフラッシュ部とを有するアルミニウム合金鍛造材であって、Mg:0.6〜1.8%、Si:0.4〜1.8%を含み、更に、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上を含み、残部Alおよび不可避的不純物からなり、前記製品部とフラッシュ部との型割り面に対し垂直方向の結晶粒径の内、前記製品部では平均結晶粒径を300 μm 以下とするとともに、前記フラッシュ部では最大の結晶粒径を400 μm 以下と各々することである。そして、このアルミニウム合金鍛造材は前記本発明製造方法で製造されることが好ましい。
【0018】
【発明の実施の形態】
図1 に、前記足回り部品などの、人工時効硬化処理後のAl合金鍛造材1 の断面図を示す通り、本発明において、Al合金鍛造材の製品部とは、Al合金鍛造材本体である製品部2 の意味である。また、フラッシュ部とは、製品部2 とフラッシュ3 との (型割り面4 に対し垂直なST方向の) 切断面5aの部位 (切断部) の意味である。  そして、本発明においては、この鍛造材の、前記製品部2 とフラッシュ部5aとの、鍛造の際の型割り面4 に対し垂直方向の結晶粒径の内、製品部2 では平均結晶粒径を300 μm 以下とするとともに、フラッシュ部5aでは最大の結晶粒径を400 μm 以下と各々規定する。
【0019】
図2 に示すように、前記足回り部品などのAl合金鍛造材1 は、通常、金型鍛造における上型7 と下型8 の金型によって、両金型の境界にできる境界面 (分割する面) である型割り面4(パーティングラインとも言う) と、上型7 と下型8 の金型との隙間から余分なAl合金を鍛造中に排出する空間であるガッタ9 と、フラッシュランドと称される一定の隙間10を設けて鍛造される。このようにして鍛造されたAl合金鍛造材1 には、上記ガッタ9 内に、必然的に、フラッシュと称されるバリ3 が生じる。このフラッシュ3 は、鍛造後、トリムライン (フラッシュ切断線)5において、製品部2 と分離切断されるが、フラッシュ3 の一部 (例えば根元部分) は残留するように切断されるため、前記足回り部品などでは、製品部2 と、型割り面4 方向に一定長さを有するフラッシュ部5aとが必然的に存在するようになる。
【0020】
一方、図1 に示すように、Al合金鍛造材1 の製品部2 の各メタルフロー (鍛流線)6は、メタルフロー6 同士の間隔が狭くなって、そのままフラッシュ3 内に流入している。しかし、このような製品部2 とフラッシュ部5aとでは、熱間鍛造時の加工率が大きく異なる。この点、通常はフラッシュ部5aの方の加工率が80% 以上と大きくなる。このため、複数回行われる熱間鍛造においては、加工率が高いフラッシュ部5aでは、最終回次の熱間鍛造の終了温度が、特に360 ℃未満の比較的低温となると、加工歪みをより加えられたフラッシュ部5aは、溶体化処理工程において、特に、加工組織が再結晶して粗大結晶粒が発生しやすくなる。
【0021】
このフラッシュ部5aやその近傍の部分、あるいは製品部に、結晶粒粗大化が生じた場合、上記ミクロ組織を制御しても、高強度化や高靱性化が果たせない。また、このフラッシュ部5aやその近傍が構造材としての使用中に、外表面となったり、前記ST方向に引張応力が付加される場合には、厳しい腐食環境との相乗効果で、この部分に応力腐食割れが発生する可能性が高くなる。
【0022】
このため、本発明では、複数回行われる熱間鍛造に際し、製品部の最終回次の熱間鍛造の終了温度を360 ℃以上として、製品部2 とフラッシュ部5aの結晶粒粗大化を防止乃至抑制する。このフラッシュ部5aの結晶粒粗大化が抑制される条件であれば、加工率が80% 未満と低くなるAl合金鍛造材1 の製品部2 の方の結晶粒粗大化は必然的により抑制される。
【0023】
そして、フラッシュ部5aの結晶粒粗大化を防止乃至抑制の目安として、フラッシュ部5aの組織における型割り面に垂直な方向(ST 方向) の結晶粒径であって、この結晶粒径の内の最大のものの大きさを400 μm 以下、好ましくは200 μm 以下に規制する。また、前記製品部2 では、製品部2 の組織における同じく型割り面に垂直な方向(ST 方向) の平均結晶粒径を300 μm 以下に規制する。
【0024】
この製品部2 の平均結晶粒径乃至フラッシュ部5aの最大結晶粒径は、測定断面の化学エッチングによって、結晶の粒界を鮮明化させ、5 〜400 倍の投影機および光学顕微鏡によって測定する。この際、材質のバラツキを考慮するため、各Al合金鍛造材毎に1 視野、20個のロッドのAl合金鍛造材の20視野観察によって行う。そして評価は、これら鍛造製品ごとの切断面の観察結果の平均値あるいは最大値によって行う。なお、結晶粒界の判別が困難な場合は、切断面を電解エッチング後、5 〜400 倍の偏光顕微鏡を用いて測定する。
【0025】
次に、本発明Al合金鍛造材乃至鍛造材用の素材における、化学成分組成について説明する。本発明のAl合金は、自動車、船舶などの輸送機材や構造材あるいは部品用として、高強度、高靱性および耐応力腐食割れ性などの高い耐久性を保証する必要がある。
【0026】
したがって、本発明Al合金鍛造材の化学成分組成は、Al−Mg−Si系のJIS 6000系Al合金の成分規格 (JIS 6101、6111、6003、6151、6061、6N01、6063など) に相当するものとして、基本的にはMg:0.6〜1.6%、Si:0.4〜1.8%を含み、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上を含む。なお、各元素量における% 表示はすべて質量% の意味である。
【0027】
しかし、JIS 6000系Al合金の各成分規格通りにならずとも、前記本発明の諸特性を阻害しない範囲で、更なる特性の向上や他の特性を付加するための、他の元素を適宜含むなどの成分組成の変更は適宜許容される。また、溶解原料スクラップなどから必然的に混入される不純物も、本発明鍛造材の品質を阻害しない範囲で許容される。
【0028】
次に、本発明Al合金鍛造材の各元素の含有量について、臨界的意義や好ましい範囲について説明する。
【0029】
Mg:0.6〜1.8%。
Mgは人工時効処理により、Siとともにβ’’相ならびにβ’ 相として析出し、最終製品使用時の高強度 (耐力) を付与するために必須の元素である。Mgの0.6%未満の含有では、人工時効処理時の時効硬化量が低下する。一方、1.8%を越えて含有されると、強度 (耐力) が高くなりすぎ、鍛造性を阻害する。また、溶体化処理後の焼き入れ途中に多量のMg2 Siや単体Siが析出しやすく、却って、強度、靱性、耐食性などを低下させる。したがって、Mgの含有量は0.6 〜1.8%の範囲とする。
【0030】
Si:0.4〜1.8%。
SiもMgとともに、人工時効処理により、β’’相ならびにβ’ 相として析出して、最終製品使用時の高強度 (耐力) を付与するために必須の元素である。Siの0.4%未満の含有では人工時効処理で十分な強度が得られない。一方、1.8%を越えて含有されると、鋳造時および溶体化処理後の焼き入れ途中で、粗大な単体Si粒子が晶出および析出して、前記した通り、耐食性と靱性を低下させる。また、過剰Siが多くなって、高耐食性と高靱性、高疲労特性を得ることができない。更に伸びが低くなるなど、加工性も阻害する。したがって、Siの含有量は0.4 〜1.8%の範囲とする。
【0031】
Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上。これらの元素は均質化熱処理時およびその後の熱間鍛造時に、Fe、Mn、Cr、Zr、Si、Alなどがその含有量に応じて選択的に結合したAl−Mn 系、Al−Cr 系、Al−Zr 系金属間化合物であり、(Fe 、Mn、Cr)SiAl12、AlZr 、(AlSi)Zr に代表される分散粒子 (分散相) を生成する。これらの分散粒子は再結晶後の粒界移動を妨げる効果があるため、前記ST方向の結晶粒の粗大化を防止するとともに、本発明Al合金鍛造材全体に渡って、微細な結晶粒や亜結晶粒を得ることができる。この結果、前記フラッシュ部のST方向の最大結晶粒を400 μm 以下、好ましくは200 μm 以下、また、前記製品部のST方向の平均結晶粒を300 μm 以下と、各々微細化させることができる。また、Mn、Cr、Zrは固溶による強度およびヤング率の増大も見込める。
【0032】
Mn、Cr、Zrの含有量が少なすぎると、これらの効果が期待できず、一方、これらの元素の過剰な含有は溶解、鋳造時に粗大な金属間化合物や晶出物を生成しやすく、破壊の起点となり、靱性や疲労特性を低下させる原因となる。このため、これらの元素は各々、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の範囲で一種または二種以上含有させる。
【0033】
Cu:0.50%以下。Cuは、Al合金鍛造材の組織の応力腐食割れや粒界腐食の感受性を著しく高め、Al合金鍛造材の耐食性や耐久性を低下させる。したがって、本発明では、この観点からCu含有量をできるだけ少なく規制する。しかし、一方で、Cuは固溶強化にて強度の向上に寄与する他、時効処理に際して、最終製品の時効硬化を著しく促進する効果も有する。なお、Cu含有量を少なくすると、高純度地金を使用する必要があり、鋳造コストがかかる問題もある。したがって、Cuは0.50% 以下の含有まで許容する。
【0034】
Fe:0.40%以下。Al合金に不純物として含まれるFeは、本発明で問題とする粗大な晶出物を生成する。これらの晶出物は、前記した通り、破壊靱性および疲労特性などを劣化させる。したがって、Feの含有量は0.40% 以下、より好ましくは0.35% 以下に規制することが好ましい。
【0035】
水素:0.25 cc/100g Al以下。水素(H)は、特に、鍛造材の加工度が小さくなる場合、水素に起因する気泡が鍛造等加工で圧着せず、破壊の起点となるため、靱性や疲労特性を著しく低下させる。そして、高強度化した輸送機の構造材などにおいては、特に水素による影響が大きい。したがって、 Al 100g当たりの水素濃度は0.25 cc 以下のできるだけ少ない含有量とすることが好ましい。
【0036】
Zn、Ti、B 、Be、V 等。  Zn、Ti、B 、Be、V 等は、各々目的に応じて、選択的に含有される元素である。
Zn:0.005〜1.0%。Znは人工時効時において、MgZn2 を微細かつ高密度に析出させ高い強度を実現させる。また、固溶したZnは粒内の電位を下げ、腐食形態を粒界からではなく、全面的な腐食として、粒界腐食や応力腐食割れを結果として軽減する効果が期待できる。しかし、Znの0.005%未満の含有では人工時効で十分な強度が得られず、前記耐食性の向上効果もない。一方、1.0%を越えて含有されると、耐蝕性が顕著に低下する。したがって、選択的に含有させる場合のZnの含有量は0.005 〜1.0%の範囲とすることが好ましい。
【0037】
Ti:0.001〜0.1%。Tiは鋳塊の結晶粒を微細化し、押出、圧延、鍛造時の加工性を向上させるために添加する元素である。しかし、Tiの0.001%未満の含有では、加工性向上の効果が得られず、一方、Tiを0.1%を越えて含有すると、粗大な晶析出物を形成し、前記加工性を低下させる。したがって、選択的に含有させる場合のTiの含有量は0.001 〜0.1%の範囲とすることが好ましい。
【0038】
B:1 〜300ppm。B はTiと同様、鋳塊の結晶粒を微細化し、押出、圧延、鍛造時の加工性を向上させるために添加する元素である。しかし、B の1ppm未満の含有では、この効果が得られず、一方、300ppmを越えて含有されると、やはり粗大な晶析出物を形成し、前記加工性を低下させる。したがって、選択的に含有させる場合のB の含有量は1 〜300ppmの範囲とすることが好ましい。
【0039】
Be:0.1〜100ppm。Beは空気中におけるAl溶湯の再酸化を防止するために含有させる元素である。しかし、0.1ppm未満の含有では、この効果が得られず、一方、100ppmを越えて含有されると、材料硬度が増大し、前記加工性を低下させる。したがって、選択的に含有させる場合のBeの含有量は0.1 〜100ppmの範囲とすることが好ましい。
【0040】
V:0.15% 以下。V は、Mn、Cr、Zrなどと同様に、均質化熱処理時およびその後の熱間鍛造時に、分散粒子 (分散相) を生成する。これらの分散粒子は再結晶後の粒界移動を妨げる効果があるため、微細な結晶粒を得ることができる。しかし過剰な含有は溶解、鋳造時に粗大なAl−Fe−Si−V系の金属間化合物や晶析出物を生成しやすく、破壊の起点となり、靱性や疲労特性を低下させる原因となる。したがって、V の含有は0.15% 以下まで許容する。
【0041】
次に、本発明におけるAl合金鍛造材の好ましい製造方法について述べる。本発明におけるAl合金鍛造材の製造自体は、前記鍛造温度を除き、常法により製造が可能である。例えば、前記Al合金成分範囲内に溶解調整されたAl合金溶湯を鋳造する場合には、例えば、連続鋳造圧延法、半連続鋳造法(DC鋳造法)、ホットトップ鋳造法等の通常の溶解鋳造法を適宜選択して鋳造する。
【0042】
なお、Al合金鋳塊の結晶粒を微細化して、鍛造材を高強度、高靱性化するためには、Al合金溶湯を、10℃/sec以上の冷却速度で鋳造して鋳塊とすることが好ましい。また、結晶粒径は150 μm 以下であることが好ましい。
【0043】
次いで、このAl合金鋳塊 (鋳造材) の均質化熱処理温度は490 〜 570℃の温度範囲とすることが好ましい。均質化熱処理温度が570 ℃を越えて高過ぎると、バーニング等が生じ、鍛造割れの原因となる。また、鍛造製品での靱性、疲労特性などの機械的な特性を低下させる。また、分散粒子が粗大化し、結晶粒微細化効果を発揮する分散粒子自体の数も不足する。
【0044】
一方、均質化熱処理温度が490 ℃未満と低過ぎると、鍛造製品を高強度化、高靱性化することが難しくなる。
【0045】
この均質化熱処理の後に、メカニカル鍛造や油圧鍛造等により熱間鍛造して、最終製品形状( ニアネットシェイプ) のAl合金鍛造材に成形する。この際、製品部の熱間鍛造開始温度を450 〜570 ℃とする。この熱間鍛造開始温度が450 ℃未満であれば、特に再加熱無しで複数回行われる熱間鍛造において、最終回次の製品部の熱間鍛造の終了温度を360 ℃以上に保証することが困難となる。また、熱間鍛造加工自体も困難となる。一方、熱間鍛造開始温度が570 ℃を越えた場合、摩擦熱により局部融解して鍛造加工割れを生じやすくなる。
【0046】
そして、鍛造後、必要な強度および靱性、耐食性を得るためのT6 (溶体化処理後、最大強さを得る人工時効硬化処理) 、T7 (溶体化処理後、最大強さを得る人工時効硬化処理条件を超えて過剰時効硬化処理) 、T8 (溶体化処理後、冷間加工を行い、更に最大強さを得る人工時効硬化処理) 等の調質処理を適宜行う。また、均質化熱処理、溶体化処理には、空気炉、誘導加熱炉、硝石炉などが適宜用いられる。更に、人工時効硬化処理には、空気炉、誘導加熱炉、オイルバスなどが適宜用いられる。
【0047】
前記溶体化処理後の焼き入れ処理の冷却は水冷が好ましい。焼き入れ処理時の冷却速度が低くなると、粒界上にMgSi 、Si等が析出し、人工時効後の製品において、粒界破壊が生じ易くなり、靱性ならびに疲労特性を低くする。また、冷却途中に、粒内にも、安定相MgSi 、Siが形成され、人工時効時に析出するβ’ 相、β’’相の析出量が減るため、強度が低下する。一方、冷却速度が高くなると、焼入歪み量が多くなり、焼入後に、矯正工程が新たに必要となったり、矯正工程の工数が増す。したがって、製品製造工程を短縮し、低コスト化するためには、焼入歪みが緩和される50〜85℃の温湯焼入が好ましい。ここで、温湯焼入温度が50℃未満では焼入歪みが大きくなり、85℃を越えると冷却速度が低くなりすぎ、靱性ならびに疲労特性、強度が低くなる。
【0048】
また、前記T7調質材では粒界上に析出するβ’ 相の割合が高くなる。このβ’ 相は腐食環境下で溶出しにくく、粒界腐食感受性を低くし、耐応力腐食割れ性を高める。一方、前記T6材で多く析出するβ’’相は腐食環境下で溶出しやすく、粒界腐食感受性を高くし、耐応力腐食割れ性を低める。したがって、Al合金鍛造材を前記T7材とすることで、耐力は若干低くなるものの、他の調質処理に比して、耐食性はより高くなる。
【0049】
なお、Al合金鍛造材に残留する鋳造組織を無くし、晶出物を破壊および微細化し、強度と靱性ならびに疲労特性をより向上させるために、Al合金鋳塊を均質化熱処理後、押出や圧延加工した後に、前記鍛造を行っても良い。
【0050】
【実施例】
次に、本発明の実施例を説明する。表1 に示す合金番号1 〜5 の化学成分組成のAl合金鋳塊 (Al合金鋳造材、いずれも直径φ82mmの丸棒) を、半連続鋳造法により、20℃/ sec の冷却速度により鋳造した。なお、この表1 に示す合金番号1 〜5 は全て本発明範囲内の化学成分組成である。また、表1 に示す合金番号1 〜5 の100gのAl中のH2 濃度は全て0.10〜0.15mlであった。
【0051】
そして、この鋳塊の外表面を厚さ3mm 面削して、長さ500mm に切断後、表2 に示すように、均質化熱処理条件 (温度と時間) 、熱間鍛造開始温度、熱間鍛造終了温度、溶体化処理温度、人工時効硬化処理条件とを種々変えて、鍛造材を作製した。均質化熱処理温度までの昇温時間は1 〜4 時間とした。また、溶体化処理は、空気炉を用いて、昇温時間を1 〜2 時間として行い、溶体化処理した後60℃の温水に焼入れを行い、その後30分以内に人工時効硬化処理を行った。なお、表2 の熱間鍛造開始温度と熱間鍛造終了温度とは、鍛造材の製品部の外表面測定温度である。
【0052】
熱間鍛造は、各例とも、前記図2 に示した上下金型を用いたメカニカル鍛造により、フラッシュランドの隙間1.5 〜3mm で、製品部の加工率が3 回の鍛造の合計で50% 、フラッシュ部の加工率が3 回の鍛造の合計で90〜95% と一定になるように、再加熱なしに3 回鍛造した。
【0053】
なお、フラッシュ部の加工率C は、鍛造材の平均結晶粒間隔A と鋳塊の平均セル層サイズB とを用い、C=[(B−A)/B] ×100%の式により算出した。鋳塊の平均セル層サイズB は鋳塊の面削前において、鋳込み方向に対する垂直面で、鋳塊外表面から中心部までを4 等分し、この鋳塊外表面から中心部への計5 箇所での平均値を用いた。一方、製品部の加工率D も、フラッシュ部の加工率C と同様に算出した。この際、加工率が小さく、明瞭なフローラインを形成しない場合には、鍛造した材料に残存する鋳塊セル層の大きさ( 最小長方向) E を用いて、D=[(B−E)/B] ×100%の式により算出した。D
【0054】
これら各鍛造材の特性を表3 に示す。なお、表2 、3 のAl合金番号は表1 のAl合金番号と対応している。この際、製品部2 の平均結晶粒径乃至フラッシュ部5aの最大結晶粒径は前記した要領で測定した。また、各鍛造材より各々引張試験片A (L方向) とシャルピー試験片B (LT 方向) を各5 個づつ採取し、引張強度(MPa) 、0.2%耐力(MPa) 、伸び(%) 、シャルピー衝撃値、等を各々測定し、各平均値を求めた。
【0055】
また、応力腐食割れ試験は、各鍛造材より、フラッシュ部を含む Cリングの試験片を採取して行った。応力腐食割れ試験条件は、前記 Cリング試験片をASTM G47の交互浸漬法の規定に準じて行った。但し、試験条件は、更に、鍛造材が切断面5a部分に対しST方向に引張応力が付加されて使用されることを模擬して、C リング試験片のST方向に、前記機械的特性の試験片のL 方向の耐力の75% の応力を負荷した状態とした。この状態で、C リング試験片の塩水への浸漬と引き上げを繰り返して90日間行い、試験片の応力腐食割れ発生の有無を確認した。これらの結果を、応力腐食割れが発生している場合を×、応力腐食割れではないが、応力腐食割れに至る可能性の高い粒界腐食が発生している場合を△、応力腐食割れや粒界腐食が発生していない場合 (表面的な全面腐食が発生している場合を含む) を○として、表2 に示す。
【0056】
表3 から明らかな通り、表1 の番号1 〜6 までの本発明範囲内の化学成分組成とし、製品部の、熱間鍛造開始温度を450 〜570 ℃とするとともに最終熱間鍛造終了温度を360 ℃以上とした鍛造材 (発明例)1と2 、4 と5 、8 と9 、12〜14、16〜21は、溶体化および焼き入れ処理と人工時効硬化処理後の鍛造材の、フラッシュ部のST方向の最大結晶粒が400 μm 以下、また、前記製品部のST方向の平均結晶粒が300 μm 以下であり、耐力 (σ0.2)が290MPa以上およびシャルピー衝撃値の平均値が15J/cm2 以上と、強度、靱性が高く、また、応力腐食割れ性にも優れている。
【0057】
これに対し、熱間鍛造開始温度が450 ℃未満乃至最終熱間鍛造終了温度が360 ℃未満である鍛造材 (比較例) 3 、6 、7 、10、11、15は、特に同じAl合金の発明例に比して、強度、靱性が発明例に比して劣り、また、応力腐食割れ性も著しく劣っている。したがって、これらの結果から、本発明熱間鍛造開始温度と最終熱間鍛造終了温度との臨界的な意義が分かる。
【0058】
【表1】
【0059】
【表2】
【0060】
【表3】
【0061】
【発明の効果】
本発明によれば、製品部とフラッシュ部とで熱間鍛造加工率が大きく異なっても、製品部とフラッシュ部ともに、結晶粒の微細化が可能で、高強度、高靱性でかつ耐食性にも優れた輸送機構造材用アルミニウム合金鍛造材の製造方法およびアルミニウム合金鍛造材を提供することができる。したがって、Al−Mg−Si系アルミニウム合金鍛造材の輸送機用への用途の拡大を図ることができる点で、多大な工業的な価値を有するものである。
【図面の簡単な説明】
【図1】本発明Al合金鍛造材で規定する組織部分のST方向断面のマクロ組織を示す断面図である。
【図2】本発明Al合金鍛造用金型を示す断面図である。
1: Al合金鍛造材、2:製品部、3:フラッシュ、4:型割り面、
5: フラッシュ切断線、6:メタルフロー、7:上型、8:下型、
9: ガッタ、10: フラッシュスタンド、
[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a method for producing an aluminum alloy forging for a transport structural material which has high strength, high toughness and excellent corrosion resistance such as stress corrosion cracking resistance (hereinafter, aluminum is simply referred to as Al) and aluminum alloy It relates to forged materials.
[0002]
[Prior art]
As is well known, AA to JIS 6000 series (Al-Mg-Si series) are used for structural materials and parts of transporting machines such as vehicles, ships, aircrafts, motorcycles and automobiles, particularly as underbody parts such as upper arms and lower arms. Forging materials such as Al alloys are used. The 6000 series Al alloy forged material has high strength, high toughness, and relatively excellent corrosion resistance. Moreover, the 6000 series Al alloy itself is also excellent in recyclability since the amount of alloying elements is small and scrap can be easily reused as a 6000 series Al alloy melting raw material.
[0003]
These 6000 series Al alloy forged materials are subjected to hot forging (die forging) such as mechanical forging and hydraulic forging after homogenizing heat treatment of the Al alloy cast material, and then to solution heat treatment, quenching and artificial aging hardening. Performed and subjected to temper treatment. Note that an extruded material obtained by once extruding a cast material may be used as the forged material.
[0004]
In recent years, structural materials for these transporters have also been required to have higher strength and higher toughness after being made thinner. For this reason, various attempts have been made to improve the microstructure of an Al alloy cast material or an Al alloy forged material.
[0005]
For example, by reducing the average particle size of crystal precipitates (crystals and precipitates) of a 6000 series Al alloy casting material to 8 μm or less and reducing the dendrite secondary arm interval (DAS) to 40 μm or less, It has been proposed to improve the strength and toughness of a forged alloy material (for example, see Patent Documents 1 and 2).
In addition, by controlling the average grain size and average interval of crystallized substances and crystal precipitates in the crystal grains and grain boundaries of the 6000 series Al alloy forged material, the Al alloy forged material has higher strength and toughness. (For example, see Patent Documents 3 and 4). These controls can enhance the corrosion resistance against intergranular corrosion and stress corrosion cracking.
[0006]
[Patent Document 1]
JP-A-07-145440
[Patent Document 2]
JP 06-256880 A
[Patent Document 3]
JP 2000-144296 A
[Patent Document 4]
JP 2001-107168 A
[0007]
However, in these 6000 series Al alloy forgings, there is a tendency that in the above-described forging and solution treatment, the processed structure is recrystallized to generate coarse crystal grains. When these coarse crystal grains are generated, even if the microstructure is controlled, high strength and high toughness cannot be achieved, and the corrosion resistance is reduced.
Conventionally, in order to suppress the generation of the coarse crystal grains, hot forging is performed at a relatively high temperature of 450 to 570 ° C. after adding a transition element having a crystal grain refining effect such as Mn, Zr, and Cr. Is known to be started (for example, see Patent Document 5).
[0008]
[Patent Document 5]
JP-A-5-247574
[0009]
[Problems to be solved by the invention]
However, even if the forging start temperature is set to a relatively high temperature of 450 to 570 ° C., the temperature of the forged material at the end of forging is not compared in hot forging where a plurality of forging steps are performed without reheating or with reheating. Extremely low temperatures are highly possible. When the temperature of the forged material at the end of the forging becomes relatively low, there is a possibility that coarse crystal grains having a recrystallized work structure may be generated, particularly in a flash portion described later.
[0010]
In addition, in a die forged product such as the undercarriage part, a product part and a flash part are inevitably present, and such a product part and the flash part have greatly different working rates during hot forging. For example, usually, in hot forging performed a plurality of times without reheating, the working ratio of the product portion is 80% or less even in a total of multiple forgings, and the working ratio of the flash portion is 80% in a total of multiple forgings. % And different. In such hot forging in which the working ratio of the product part and the flash part are different, even in the case where the hot forging starting temperature is a relatively high temperature of 450 to 570 ° C., the flash part to which the working strain is further applied has a solution solution. In the processing step, particularly, the processed structure is likely to recrystallize to generate coarse crystal grains.
[0011]
On the other hand, when trying to reduce the processing rate of the flash part in order to reduce the processing distortion of the flash part, the processing rate of the product part necessarily decreases, the cast structure remains, and the strength and toughness decrease. There is also a possibility that the shape accuracy may not be obtained.
[0012]
This problem can also occur in Japanese Patent Application No. 2002-76567 previously proposed by the present inventors. That is, in the present invention, Mg: 0.6-1.8%, Si: 0.4-1.8%, Mn: 0.01-0.6%, Cr: 0.1-0. .2% and Zr: an aluminum alloy forged material containing one or more of 0.1 to 0.2%, the balance being Al and unavoidable impurities, the aluminum alloy forged material after artificial aging treatment, 0.2% proof stress of 300MPa or more and Charpy impact value of 10J / cm 2 Further, it is assumed that, in the structure of the cut surface between the product and the flash, the crystal grain size in the direction perpendicular to the mold surface, and the largest one of the crystal grain sizes is 400 μm or less. Features.
[0013]
However, also in the present invention, the preferred hot forging starting temperature is a relatively low temperature of 350 to 450 ° C. in all the examples, and in the case of hot forging in which a plurality of forging steps are performed without reheating, It is quite possible that the temperature of the forged material at the end of forging will be relatively low. As a result, there is still a possibility that coarse crystal grains in which the processed structure is recrystallized are generated in the flash part and further in the product part.
[0014]
In view of such circumstances, an object of the present invention is to provide a transport structural material capable of refining crystal grains in both a product part and a flash part even when a hot forging rate is greatly different between the product part and the flash part. It is an object of the present invention to provide a method for manufacturing an aluminum alloy forging for use and an aluminum alloy forging.
[0015]
[Means for Solving the Problems]
In order to achieve this object, the gist of the method for manufacturing an aluminum alloy forging for a transport aircraft structural material according to the present invention is a method for manufacturing an aluminum alloy forging having a product part and a flash part, wherein Mg: 0.6 -1.8%, Si: 0.4-1.8%, Mn: 0.01-0.9%, Cr: 0.01-0.25%, and Zr: 0.01-0. When hot forging is performed on an Al-Mg-Si-based aluminum alloy casting material containing one or more than 20% of aluminum alloy, the hot forging starting temperature of the product part is set to 450 to 570 ° C. At the same time, the final hot forging end temperature of the product part is set to 360 ° C. or higher, and the forged material after solution heat treatment, quenching, and artificial age hardening is perpendicular to the parting surface between the product part and the flash part. Of the crystal grain size, The crystal grain size with a 300 [mu] m or less, is to each and 400 [mu] m or less maximum crystal grain size in the flash unit.
[0016]
In the present invention, the product part and the flash part are such that the processing rate of the product part is 80% or less in total of a plurality of forgings, and the processing rate of the flash part is 80% or more in total of a plurality of forgings. In hot forging with different working ratios, the hot forging start temperature of the product part is set to 450 to 570 ° C., and the end temperature of the final round hot forging of the product part is set to 360 ° C. or more. By defining the start temperature and the end temperature of the final round hot forging of the product part in this way, it is possible to prevent the crystal grain size of the product part from becoming coarse and to reduce the crystal grain size of the flash part. Even coarsening can be prevented or suppressed. In the present invention, the end temperature and the start temperature of the hot forging, or the temperature specified in the homogenization heat treatment, the tempering treatment after the forging, etc. are all the temperatures of the outer surface of the cast material or the forged product part. .
For example, even if the start temperature of hot forging of the product part is 450 to 570 ° C., if the end temperature of hot forging of the product part is lower than 360 ° C., fine grains of Mn, Zr, Cr, etc. Even if a transition element having a quenching effect is added, the average grain size in the product part of the forged material after tempering, of the crystal grain diameter in the direction perpendicular to the parting plane of the product part and the flash part, Is not more than 300 μm, and the maximum crystal grain size in the flash part cannot be made not more than 400 μm. As a result, it becomes impossible to increase the strength and toughness of the aluminum alloy forged material, and in particular, it is not possible to increase the corrosion resistance of the flash portion.
[0017]
In order to achieve the above object, the gist of the present invention for an aluminum alloy forged material for a transport aircraft structural material is an aluminum alloy forged material having a product part and a flash part, and Mg: 0.6 to 1.8. %, Si: 0.4 to 1.8%, and further, Mn: 0.01 to 0.9%, Cr: 0.01 to 0.25%, and Zr: 0.01 to 0.20%. One or more of them, the balance being Al and unavoidable impurities, of which the product part has an average crystal grain size of 300 μm in the direction perpendicular to the parting plane of the product part and the flash part. The maximum crystal grain size in the flash portion is set to 400 μm or less. And this aluminum alloy forging material is preferably manufactured by the manufacturing method of the present invention.
[0018]
BEST MODE FOR CARRYING OUT THE INVENTION
FIG. 1 shows a cross-sectional view of an Al alloy forged material 1 after artificial age hardening treatment, such as the underbody part, in the present invention, the product part of the Al alloy forged material is the Al alloy forged material body. Product part 2 The flash portion means a portion (cut portion) of the cut surface 5a (in the ST direction perpendicular to the mold surface 4) between the product portion 2 and the flash 3. In the present invention, of the forged material, the average grain size of the product portion 2 and the flash portion 5a is the average grain size of the product portion 2 in the direction perpendicular to the parting surface 4 at the time of forging. Is set to 300 μm or less, and the maximum crystal grain size in the flash portion 5a is specified to be 400 μm or less.
[0019]
As shown in FIG. 2, the Al alloy forging material 1 such as the underbody part is usually formed by an upper die 7 and a lower die 8 in die forging. A mold surface 4 (also referred to as a parting line), a gutter 9 which is a space for discharging excess Al alloy from the gap between the upper mold 7 and the lower mold 8 during forging, and a flash land. Forging is performed by providing a constant gap 10 referred to as “forging”. In the forged Al alloy material 1 thus forged, burrs 3 called flashes are inevitably generated in the gutter 9. After the forging, the flash 3 is cut off from the product part 2 at a trim line (flash cutting line) 5, but a part of the flash 3 (for example, a root portion) is cut so as to remain. In the peripheral parts, the product part 2 and the flash part 5a having a certain length in the direction of the mold parting plane 4 necessarily exist.
[0020]
On the other hand, as shown in FIG. 1, each metal flow (forging flow line) 6 of the product part 2 of the Al alloy forged material 1 flows into the flash 3 as it is because the interval between the metal flows 6 becomes narrow. . However, the working ratio at the time of hot forging differs greatly between the product part 2 and the flash part 5a. In this regard, the processing rate of the flash portion 5a is usually as large as 80% or more. For this reason, in the hot forging performed a plurality of times, in the flash portion 5a having a high working ratio, when the end temperature of the hot forging in the final round becomes a relatively low temperature, particularly less than 360 ° C., a working strain is added. In the solution treatment step, particularly, the processed flash structure of the flash portion 5a is likely to recrystallize to generate coarse crystal grains.
[0021]
If crystal grains are coarsened in the flash part 5a, in the vicinity thereof, or in the product part, high strength and high toughness cannot be achieved even if the microstructure is controlled. If the flash portion 5a or its vicinity becomes an outer surface or a tensile stress is applied in the ST direction during use as a structural material, a synergistic effect with a severe corrosive environment may occur in this portion. The possibility of occurrence of stress corrosion cracking increases.
[0022]
For this reason, in the present invention, in the hot forging performed a plurality of times, the end temperature of the final hot forging of the product part is set to 360 ° C. or more to prevent crystal grain coarsening of the product part 2 and the flash part 5a. Suppress. Under the condition that the coarsening of the crystal grains in the flash portion 5a is suppressed, the coarsening of the crystal grains in the product part 2 of the Al alloy forged material 1 in which the working ratio is reduced to less than 80% is necessarily further suppressed. .
[0023]
The grain size of the structure of the flash portion 5a in the direction (ST direction) perpendicular to the parting plane is used as a standard for preventing or suppressing the coarsening of the crystal grain of the flash portion 5a. The maximum size is restricted to 400 μm or less, preferably 200 μm or less. In the product part 2, the average crystal grain size of the structure of the product part 2 in the direction (ST direction) also perpendicular to the die surface is regulated to 300 μm or less.
[0024]
The average crystal grain size of the product part 2 to the maximum crystal grain size of the flash part 5a are measured by a projector and an optical microscope of 5 to 400 times magnification by sharpening the crystal grain boundaries by chemical etching of the measurement section. At this time, in order to consider the variation in the material, the observation is performed in one visual field for each Al alloy forging material, and in 20 visual fields of the Al alloy forging material of 20 rods. The evaluation is performed based on the average value or the maximum value of the observation results of the cut surface for each forged product. If it is difficult to determine the crystal grain boundaries, the cut surface is electrolytically etched and then measured using a 5- to 400-fold polarizing microscope.
[0025]
Next, the chemical composition of the Al alloy forged material of the present invention or the material for the forged material will be described. The Al alloy of the present invention is required to guarantee high durability such as high strength, high toughness and stress corrosion cracking resistance for use in transportation equipment, structural materials and parts for automobiles and ships.
[0026]
Therefore, the chemical composition of the forged aluminum alloy according to the present invention corresponds to the component standard (JIS 6101, 6111, 6003, 6151, 6061, 6N01, 6063, etc.) of the Al-Mg-Si JIS 6000 series Al alloy. Basically, Mg: 0.6-1.6%, Si: 0.4-1.8%, Mn: 0.01-0.9%, Cr: 0.01-0.25 % And Zr: 0.01 to 0.20%. In addition,% display in each element amount means mass%.
[0027]
However, other elements for further improving the properties and adding other properties are appropriately included within a range that does not impair the above-mentioned various properties of the present invention, even if the component specifications of the JIS 6000 series Al alloy are not met. Changes in the component composition, such as, are appropriately permitted. Further, impurities inevitably mixed in from the raw material scrap and the like are allowed as long as the quality of the forged material of the present invention is not impaired.
[0028]
Next, regarding the content of each element of the aluminum alloy forged material of the present invention, critical significance and preferred ranges will be described.
[0029]
Mg: 0.6-1.8%.
Mg is an essential element that precipitates together with Si as a β ″ phase and a β ′ phase by artificial aging treatment and imparts high strength (proof stress) when the final product is used. When the content of Mg is less than 0.6%, the age hardening amount during the artificial aging treatment decreases. On the other hand, if the content exceeds 1.8%, the strength (proof stress) becomes too high, and the forgeability is impaired. Also, during the quenching after the solution treatment, a large amount of Mg 2 Si or elemental Si is likely to precipitate, but rather decreases strength, toughness, corrosion resistance, and the like. Therefore, the content of Mg is set in the range of 0.6 to 1.8%.
[0030]
Si: 0.4 to 1.8%.
Si is also an element that precipitates together with Mg as a β ″ phase and a β ′ phase by artificial aging to impart high strength (proof stress) when the final product is used. If the content of Si is less than 0.4%, sufficient strength cannot be obtained by the artificial aging treatment. On the other hand, when the content exceeds 1.8%, coarse single Si particles are crystallized and precipitated during casting and during quenching after solution treatment, and as described above, the corrosion resistance and toughness are reduced. . In addition, excess Si increases, and high corrosion resistance, high toughness, and high fatigue characteristics cannot be obtained. Further, workability is impaired, such as lower elongation. Therefore, the content of Si is set in the range of 0.4 to 1.8%.
[0031]
One or more of Mn: 0.01 to 0.9%, Cr: 0.01 to 0.25% and Zr: 0.01 to 0.20%. At the time of the homogenizing heat treatment and the subsequent hot forging, these elements are Al-Mn-based, Al-Cr-based, in which Fe, Mn, Cr, Zr, Si, Al, and the like are selectively bonded according to their contents. Al-Zr based intermetallic compound, (Fe 2, Mn, Cr) 3 SiAl 12 , Al 3 Zr, (AlSi) 3 A dispersed particle (dispersed phase) represented by Zr is generated. Since these dispersed particles have the effect of hindering the movement of the grain boundary after recrystallization, the crystal grains in the ST direction are prevented from becoming coarse, and the fine grains and sub-crystals are formed throughout the Al alloy forged material of the present invention. Crystal grains can be obtained. As a result, the maximum crystal grain in the ST direction of the flash part can be reduced to 400 μm or less, preferably 200 μm or less, and the average crystal grain in the ST direction of the product part can be reduced to 300 μm or less. In addition, Mn, Cr, and Zr can be expected to increase in strength and Young's modulus due to solid solution.
[0032]
If the contents of Mn, Cr, and Zr are too small, these effects cannot be expected. On the other hand, excessive contents of these elements are likely to produce coarse intermetallic compounds and crystallized substances during melting and casting, resulting in destruction. , Which causes toughness and fatigue characteristics to be reduced. Therefore, each of these elements contains one or more of Mn: 0.01 to 0.9%, Cr: 0.01 to 0.25%, and Zr: 0.01 to 0.20%. Let it.
[0033]
Cu: 0.50% or less. Cu remarkably increases the susceptibility of the structure of the Al alloy forging to stress corrosion cracking and intergranular corrosion, and lowers the corrosion resistance and durability of the Al alloy forging. Therefore, in the present invention, the Cu content is regulated as small as possible from this viewpoint. However, on the other hand, Cu contributes to improvement of strength by solid solution strengthening, and also has an effect of remarkably promoting age hardening of a final product during aging treatment. In addition, when the Cu content is reduced, it is necessary to use a high-purity ingot, and there is a problem that the casting cost is high. Therefore, Cu is allowed up to a content of 0.50% or less.
[0034]
Fe: 0.40% or less. Fe contained as an impurity in the Al alloy generates a coarse crystallized substance which is a problem in the present invention. These crystals degrade the fracture toughness and fatigue properties as described above. Therefore, the content of Fe is preferably regulated to 0.40% or less, more preferably 0.35% or less.
[0035]
Hydrogen: 0.25 cc / 100 g Al or less. Hydrogen (H 2 In the case of (2), particularly when the degree of processing of the forged material is small, the bubbles caused by hydrogen are not pressed by forging or the like and serve as starting points for destruction, so that the toughness and the fatigue properties are significantly reduced. And, in the structural material of a transport machine with high strength, the influence of hydrogen is particularly large. Therefore, the hydrogen concentration per 100 g of Al is preferably as low as 0.25 cc or less.
[0036]
Zn, Ti, B, Be, V and the like. Zn, Ti, B, Be, V and the like are elements that are selectively contained depending on the purpose.
Zn: 0.005 to 1.0%. Zn is MgZn during artificial aging. 2 To achieve high strength by precipitating fine and high density. In addition, the solid solution Zn lowers the potential in the grains, and can be expected to have the effect of reducing the intergranular corrosion and stress corrosion cracking as a result of not changing the corrosion form from the grain boundaries but as the entire corrosion. However, if the content of Zn is less than 0.005%, sufficient strength cannot be obtained by artificial aging, and there is no effect of improving the corrosion resistance. On the other hand, when the content exceeds 1.0%, the corrosion resistance is significantly reduced. Therefore, the content of Zn when it is selectively contained is preferably in the range of 0.005 to 1.0%.
[0037]
Ti: 0.001 to 0.1%. Ti is an element added for refining the crystal grains of the ingot and improving workability during extrusion, rolling, and forging. However, if the content of Ti is less than 0.001%, the effect of improving the workability cannot be obtained. On the other hand, if the content of Ti exceeds 0.1%, coarse crystal precipitates are formed and the workability is reduced. Lower. Therefore, the content of Ti when it is selectively contained is preferably in the range of 0.001 to 0.1%.
[0038]
B: 1 to 300 ppm. B is an element added to refine the crystal grains of the ingot and improve workability during extrusion, rolling, and forging, like Ti. However, when the content of B 2 is less than 1 ppm, this effect cannot be obtained. On the other hand, when the content of B 2 exceeds 300 ppm, coarse crystal precipitates are also formed and the processability is reduced. Therefore, the content of B 2 when it is selectively contained is preferably in the range of 1 to 300 ppm.
[0039]
Be: 0.1 to 100 ppm. Be is an element contained to prevent re-oxidation of the Al melt in the air. However, if the content is less than 0.1 ppm, this effect cannot be obtained. On the other hand, if the content is more than 100 ppm, the material hardness increases and the workability decreases. Therefore, the content of Be when selectively contained is preferably in the range of 0.1 to 100 ppm.
[0040]
V: 0.15% or less. V 2, like Mn, Cr, Zr, etc., generates dispersed particles (dispersed phase) during the homogenizing heat treatment and during the subsequent hot forging. Since these dispersed particles have an effect of hindering the movement of the grain boundary after recrystallization, fine crystal grains can be obtained. However, an excessive content tends to generate a coarse Al-Fe-Si-V-based intermetallic compound or crystal precipitate during melting and casting, becomes a starting point of fracture, and lowers toughness and fatigue properties. Therefore, the content of V is allowed up to 0.15% or less.
[0041]
Next, a preferred method of manufacturing the forged Al alloy according to the present invention will be described. Except for the forging temperature, the production itself of the Al alloy forging material in the present invention can be produced by a conventional method. For example, in the case of casting an Al alloy melt that has been melt-adjusted within the above-mentioned range of the Al alloy component, for example, ordinary melt casting such as continuous casting rolling, semi-continuous casting (DC casting), and hot-top casting is used. Casting is performed by appropriately selecting a method.
[0042]
In order to refine the crystal grains of the Al alloy ingot to increase the strength and toughness of the forged material, the molten Al alloy is cast at a cooling rate of 10 ° C./sec or more to form an ingot. Is preferred. Further, the crystal grain size is preferably 150 μm or less.
[0043]
Next, it is preferable that the homogenizing heat treatment temperature of the Al alloy ingot (cast material) be in a temperature range of 490 to 570 ° C. If the temperature of the homogenizing heat treatment is too high, exceeding 570 ° C., burning or the like occurs, which causes forging cracks. Also, it reduces mechanical properties such as toughness and fatigue properties of the forged product. In addition, the dispersed particles are coarsened, and the number of the dispersed particles that exhibit the crystal grain refining effect is insufficient.
[0044]
On the other hand, if the homogenizing heat treatment temperature is too low, less than 490 ° C., it becomes difficult to increase the strength and toughness of the forged product.
[0045]
After this homogenization heat treatment, the forging is performed by hot forging by mechanical forging, hydraulic forging, or the like to form an Al alloy forged material having a final product shape (near net shape). At this time, the hot forging start temperature of the product part is set to 450 to 570 ° C. If the hot forging start temperature is lower than 450 ° C., particularly in the hot forging performed a plurality of times without reheating, the end temperature of the hot forging of the product part in the final round can be guaranteed to be 360 ° C. or more. It will be difficult. In addition, the hot forging itself becomes difficult. On the other hand, if the hot forging start temperature exceeds 570 ° C., local melting occurs due to frictional heat, and forging cracks are likely to occur.
[0046]
Then, after forging, T6 (artificial age hardening treatment for obtaining maximum strength after solution treatment) and T7 (artificial age hardening treatment for obtaining maximum strength after solution treatment) to obtain necessary strength, toughness, and corrosion resistance. Temper treatments such as excessive age hardening beyond the conditions) and T8 (artificial age hardening to perform the cold working after the solution treatment and further obtain the maximum strength) are appropriately performed. For the homogenizing heat treatment and the solution treatment, an air furnace, an induction heating furnace, a nitrite furnace and the like are appropriately used. Further, an air furnace, an induction heating furnace, an oil bath, or the like is appropriately used for the artificial age hardening treatment.
[0047]
The cooling in the quenching treatment after the solution treatment is preferably water cooling. When the cooling rate during the quenching process decreases, Mg 2 Si, Si, and the like are precipitated, and in the product after artificial aging, grain boundary fracture is likely to occur, and the toughness and fatigue characteristics are lowered. During the cooling, the stable phase Mg 2 Si, Si is formed and β precipitated during artificial aging ' Phase, β '' Since the amount of phase precipitation is reduced, the strength is reduced. On the other hand, when the cooling rate increases, the amount of quenching distortion increases, and after quenching, a new straightening step is required or the number of steps in the straightening step is increased. Therefore, in order to shorten the product manufacturing process and reduce the cost, it is preferable to use hot water quenching at 50 to 85 ° C. in which quenching distortion is reduced. Here, if the hot water quenching temperature is less than 50 ° C., the quenching strain increases, and if it exceeds 85 ° C., the cooling rate becomes too low, and the toughness, fatigue characteristics, and strength decrease.
[0048]
In the T7 tempered material, β precipitated on the grain boundary ' The proportion of phases is higher. This β ' The phase hardly elutes in a corrosive environment, lowers intergranular corrosion susceptibility, and enhances stress corrosion cracking resistance. On the other hand, β which precipitates much in the T6 material '' The phase is easily eluted in a corrosive environment, increases intergranular corrosion susceptibility, and reduces stress corrosion cracking resistance. Therefore, when the forged Al alloy is made of the T7 material, the proof strength is slightly lowered, but the corrosion resistance is higher than other tempering treatments.
[0049]
In order to eliminate the cast structure remaining in the Al alloy forging, break and refine the crystallized material, and further improve the strength, toughness and fatigue properties, the Al alloy ingot was homogenized, extruded and rolled. After that, the forging may be performed.
[0050]
【Example】
Next, examples of the present invention will be described. An Al alloy ingot (Al alloy cast material, a round bar having a diameter of 82 mm) having an alloy composition of alloy numbers 1 to 5 shown in Table 1 was cast by a semi-continuous casting method at a cooling rate of 20 ° C / sec. . The alloy numbers 1 to 5 shown in Table 1 are all chemical component compositions within the scope of the present invention. Also, H in 100 g of Al of alloy numbers 1 to 5 shown in Table 1 2 All concentrations were 0.10-0.15 ml.
[0051]
Then, the outer surface of the ingot was chamfered to a thickness of 3 mm and cut to a length of 500 mm. Then, as shown in Table 2, homogenization heat treatment conditions (temperature and time), hot forging start temperature, hot forging Forgings were produced by variously changing the termination temperature, the solution treatment temperature, and the conditions of the artificial aging hardening treatment. The heating time up to the homogenizing heat treatment temperature was 1 to 4 hours. The solution treatment was performed using an air furnace with a temperature rise time of 1 to 2 hours. After the solution treatment, quenching was performed in hot water at 60 ° C., and then artificial aging hardening was performed within 30 minutes. . The hot forging start temperature and hot forging end temperature in Table 2 are the measured outer surface temperatures of the product part of the forged material.
[0052]
In each case, the hot forging is performed by mechanical forging using the upper and lower dies shown in FIG. 2 described above. Forging was performed three times without reheating so that the working ratio of the flash portion was constant at 90 to 95% in total of the three forgings.
[0053]
In addition, the working ratio C of the flash part was calculated using the average crystal grain spacing A of the forged material and the average cell layer size B of the ingot by the formula of C = [(BA) / B] × 100%. . The average cell layer size B of the ingot is a surface perpendicular to the casting direction, which is divided into four equal parts from the outer surface of the ingot to the center before the ingot is chamfered. The average value at each point was used. On the other hand, the processing rate D 1 of the product part was also calculated in the same manner as the processing rate C of the flash part. At this time, when the processing rate is small and a clear flow line is not formed, D = [(BE) using the size (minimum length direction) E of the ingot cell layer remaining in the forged material. / B] × 100%. D
[0054]
Table 3 shows the characteristics of these forged materials. The Al alloy numbers in Tables 2 and 3 correspond to the Al alloy numbers in Table 1. At this time, the average crystal grain size of the product part 2 to the maximum crystal grain size of the flash part 5a were measured as described above. Further, five tensile test pieces A (L direction) and five Charpy test pieces B (LT direction) were taken from each forged material, and tensile strength (MPa), 0.2% proof stress (MPa), elongation (%) ), Charpy impact value, etc. were measured, and the respective average values were determined.
[0055]
In addition, the stress corrosion cracking test was performed by collecting a C ring test piece including a flash portion from each forged material. The stress corrosion cracking test conditions were performed on the C-ring test piece in accordance with the provisions of the alternate immersion method of ASTM G47. However, the test conditions further simulated that the forged material was used by applying a tensile stress to the cut surface 5a in the ST direction in the ST direction. A state in which a stress of 75% of the proof stress in the L direction of the piece was applied. In this state, immersion and lifting of the C-ring test piece in salt water were repeatedly performed for 90 days, and the presence or absence of stress corrosion cracking of the test piece was confirmed. These results are shown as × for cases where stress corrosion cracking has occurred, and △ for cases where intergranular corrosion that is not stress corrosion cracking but is likely to lead to stress corrosion cracking has occurred. Table 2 shows the case where interfacial corrosion did not occur (including the case where superficial general corrosion occurred).
[0056]
As is clear from Table 3, the chemical composition within the range of the present invention from No. 1 to 6 in Table 1 is set, the hot forging start temperature of the product section is set to 450 to 570 ° C., and the final hot forging end temperature is set. Forged material of 360 ° C. or higher (Invention example) 1 and 2, 4 and 5, 8 and 9, 12 to 14 and 16 to 21 are flashes of forged material after solution treatment and quenching treatment and artificial age hardening treatment. The maximum crystal grain in the ST direction of the part is 400 μm or less, and the average crystal grain in the ST direction of the product part is 300 μm or less. 0.2 ) Is 290 MPa or more and the average value of the Charpy impact value is 15 J / cm. 2 As described above, the strength and toughness are high, and the stress corrosion cracking resistance is also excellent.
[0057]
On the other hand, forged materials having a hot forging start temperature of less than 450 ° C. to a final hot forging end temperature of less than 360 ° C. (Comparative Examples) 3, 6, 7, 10, 11, and 15 are particularly the same Al alloys. The strength and toughness are inferior to the invention examples, and the stress corrosion cracking properties are also extremely inferior to the invention examples. Therefore, these results show the critical significance of the hot forging start temperature and the final hot forging end temperature of the present invention.
[0058]
[Table 1]
[0059]
[Table 2]
[0060]
[Table 3]
[0061]
【The invention's effect】
According to the present invention, even if the hot forging ratio is significantly different between the product part and the flash part, the product part and the flash part can have fine crystal grains, and have high strength, high toughness, and corrosion resistance. An excellent method for producing an aluminum alloy forging for a transport aircraft structural material and an aluminum alloy forging can be provided. Therefore, the use of the forged Al-Mg-Si-based aluminum alloy can be expanded for use in transport vehicles, and thus has great industrial value.
[Brief description of the drawings]
FIG. 1 is a cross-sectional view showing a macro structure of a cross section in the ST direction of a structure portion defined by an aluminum alloy forged material of the present invention.
FIG. 2 is a sectional view showing a die for forging an Al alloy of the present invention.
1: Forged Al alloy, 2: Product part, 3: Flash, 4: Molded surface,
5: Flash cutting line, 6: Metal flow, 7: Upper die, 8: Lower die,
9: Gutta, 10: Flash stand,

Claims (6)

  1. 製品部とフラッシュ部とを有するアルミニウム合金鍛造材の製造方法であって、Mg:0.6〜1.8%、Si:0.4〜1.8%を含み、更に、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上を含むAl−Mg−Si系アルミニウム合金鋳造材を、均質化熱処理後に熱間鍛造を行うに際し、前記製品部の熱間鍛造開始温度を450 〜570 ℃とするとともに、製品部の最終熱間鍛造終了温度を360 ℃以上とし、溶体化および焼き入れ処理と人工時効硬化処理後の鍛造材の、前記製品部とフラッシュ部との型割り面に対し垂直方向の結晶粒径の内、前記製品部では平均結晶粒径を300 μm 以下とするとともに、前記フラッシュ部では最大の結晶粒径を400 μm 以下と各々することを特徴とする輸送機構造材用アルミニウム合金鍛造材の製造方法。A method for producing an aluminum alloy forging having a product part and a flash part, comprising: 0.6 to 1.8% of Mg, 0.4 to 1.8% of Si, and further Mn: 0.01 -0.9%, Cr: 0.01% -0.25%} and Zr: 0.01% -0.20%}. When hot forging is performed later, the hot forging start temperature of the product part is set to 450 ° C. to 570 ° C., and the final hot forging end temperature of the product part is set to 360 ° C. or more. Of the forged material after the hardening treatment, of the crystal grain diameter in the direction perpendicular to the parting plane between the product part and the flash part, the product part has an average crystal grain diameter of 300 μm or less, and the flash part has Largest grain The 400 [mu] m or less and each method for producing a transport aircraft structural materials for the aluminum alloy forging, characterized by.
  2. 前記人工時効硬化処理後のアルミニウム合金鍛造材製品部の0.2%耐力が290MPa以上およびシャルピー衝撃値が15J/cm2 以上である請求項1に記載の輸送機構造材用アルミニウム合金鍛造材の製造方法。 2. The aluminum alloy forged material for a transport machine structural material according to claim 1, wherein the aluminum alloy forged material product part after the artificial age hardening treatment has a 0.2% proof stress of 290 MPa or more and a Charpy impact value of 15 J / cm 2 or more. Production method.
  3. 前記均質化熱処理温度を490 〜 570℃とする請求項1または2に記載の輸送機構造材用アルミニウム合金鍛造材の製造方法。The method for producing an aluminum alloy forged material for a transport structure according to claim 1 or 2, wherein the homogenizing heat treatment temperature is 490 ° C to 570 ° C.
  4. 前記輸送機構造材が自動車足回り部品用である請求項1乃至3の何れか1項に記載の輸送機構造材用アルミニウム合金鍛造材の製造方法。The method for producing an aluminum alloy forging for a transport structure according to any one of claims 1 to 3, wherein the transport structure is for underbody parts of an automobile.
  5. 製品部とフラッシュ部とを有するアルミニウム合金鍛造材であって、Mg:0.6〜1.8%、Si:0.4〜1.8%を含み、更に、Mn:0.01 〜0.9%、Cr:0.01 〜0.25% およびZr:0.01 〜0.20% の一種または二種以上を含み、残部Alおよび不可避的不純物からなり、前記製品部とフラッシュ部との型割り面に対し垂直方向の結晶粒径の内、前記製品部では平均結晶粒径を300 μm 以下とするとともに、前記フラッシュ部では最大の結晶粒径を400 μm 以下と各々することを特徴とする輸送機構造材用アルミニウム合金鍛造材。An aluminum alloy forging material having a product part and a flash part, which contains 0.6 to 1.8% of Mg and 0.4 to 1.8% of Si, and further has a Mn of 0.01 to 0.1%. 9%, Cr: 0.01% to 0.25%} and Zr: 0.01% to 0.20%}, and the balance is composed of Al and inevitable impurities. Among the crystal grain diameters perpendicular to the parting plane, the product part has an average crystal grain diameter of 300 μm or less, and the flash part has a maximum crystal grain diameter of 400 μm or less. Alloy forgings for transport aircraft structural materials.
  6. 請求項1乃至4のいずれか1項の製造方法によって製造された請求項5に記載の輸送機構造材用アルミニウム合金鍛造材。The aluminum alloy forged material for a transport structure according to claim 5, which is manufactured by the manufacturing method according to any one of claims 1 to 4.
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