JP3657738B2

JP3657738B2 - Method for producing aluminum alloy plate for can body with low ear rate

Info

Publication number: JP3657738B2
Application number: JP12073997A
Authority: JP
Inventors: 幸男浦吉; 了東海林
Original assignee: Furukawa Sky Aluminum Corp
Current assignee: Furukawa Sky Aluminum Corp
Priority date: 1997-05-12
Filing date: 1997-05-12
Publication date: 2005-06-08
Anticipated expiration: 2017-05-12
Also published as: JPH10310837A

Description

【０００１】
【発明の属する技術分野】
本発明は、特に飲料缶胴材に適した、高強度で、しごき加工性、塗装焼付け後のフランジ成形性に優れた、耳率の低いキャンボディ用アルミニウム合金板の製造方法に関する。
【０００２】
【従来の技術】
従来、飲料缶胴材は、ＪＩＳ−３００４合金鋳塊を均質化処理し、その後、熱間圧延、冷間圧延、焼鈍、冷間圧延して製造されている。冷間圧延後、必要に応じ、仕上焼鈍、脱脂、洗浄、カッピング用潤滑油塗布などが施される。
【０００３】
飲料缶用胴材をカップ（円筒）状に絞るとカップの周縁部に凹凸が生じる場合がある。前記凸部と凹部高さのカップ高さに対する比率を耳率と言うが、前記耳率が高いと、カップ成形およびしごき成形時に耳先端からチップが飛び込んでピンホールやティアーオフが発生したり、フランジ成形後の缶の寸法精度が低下したりする。そこで耳率が高い場合は、缶ボディ成形後のトリミング量を増やすが、トリミング後も凹部が残ってしまうという問題がある。また最近増えだした缶径の小さい缶の場合、耳率の高い材料を用いると、フランジングやネッキング加工の際、フランジ長さのばらつきが大きくなり、蓋との巻締めに支障をきたすという新たな問題が生じている。このようなことから、飲料缶胴材には、これまで以上に低耳率の材料が要求されるようになってきている。
【０００４】
ところで、耳は、圧延材の結晶学的異方性に起因して生じるものであり、その高低は、熱延終了後或いは焼鈍中に進行する再結晶により形成される立方体方位の再結晶粒の集合組織成分（主に０°−９０°耳）と、圧延加工（冷間圧延）により形成される圧延集合組織成分（４５°耳）とのバランスによって決まる。
たとえば、缶強度を重視する場合は、冷間加工を高圧下率で行うため圧延集合組織が強く形成される。そこで、この場合は、熱間圧延または焼鈍条件を厳密に規定して立方体方位再結晶粒を優先成長させて対処している（特開平4-228551号、特開平6-158244号）。
しかし、近年は、缶径の縮小に伴って、耳率に対するユーザーの要求が益々厳しくなり、熱間圧延条件または焼鈍条件を規定するだけではユーザーの要求する低耳率は実現できず、耳率をさらに低くするには、熱間圧延条件に加え、その上工程の均質化処理条件などについても広く検討する必要がでてきた。
【０００５】
【発明が解決しようとする課題】
本発明は、縮径管に対しても、耳率を十分安定して低くでき、かつ成形性と強度に優れたキャンボディ用アルミニウム合金板の製造方法の提供を目的とする。
【０００６】
【課題を解決するための手段】
本発明は、Ｍｇを０．８〜１．４wt％、Ｍｎを０．７〜１．３wt％、Ｆｅを０．２〜０．５wt％、Ｓｉを０．１〜０．５wt％、Ｃｕを０．１〜０．３wt％、Ｔｉ０．００５〜０．０５wt％を単独で或いはＢ０．０００１〜０．０１wt％とともに含有し、残部がＡｌと不可避的不純物からなるアルミニウム合金鋳塊に、５６０〜６２０℃の温度範囲で１時間以上の均質化処理を施し、次いで均質化処理温度から２０℃／時間以上の冷却速度で４５０〜５５０℃まで冷却して熱間粗圧延を施すか、或いは均質化処理後そのまま室温まで冷却したのち３０℃／時間以上の昇温速度で４５０〜５５０℃まで再加熱して熱間粗圧延を施し、前記熱間粗圧延を終了板厚が１２〜５０ｍｍ、終了温度が３００〜４５０℃、最終パス圧下率ＲがＲ≦７０−０．２Ｓ（Ｓ：圧延速度ｍ／分）の条件で施し、熱間粗圧延終了後ｔ秒（ｔ＝２．８×１０⁴exp（−０．０１２Ｔ），Ｔ：熱間粗圧延終了温度℃) 以内に熱間仕上圧延を開始し、前記熱間仕上圧延をスタンド数３以上のタンデム式熱間仕上圧延機を用い、総圧下率８０％以上、各スタンドでの圧下率３０％以上、終了板厚１．６〜３．０ｍｍ、終了温度２９０℃以上の条件で施し、熱間仕上圧延後室温まで冷却し、続いて箱型焼鈍炉を用いて３００〜４５０℃で３０分以上保持して焼鈍するか、連続焼鈍炉を用いて１００℃／分以上の昇温速度で３６０〜５６０℃の温度に保持して焼鈍し、前記温度に到達後直ちに或いは１２０秒以下の時間保持後１００℃／分以上の冷却速度で７０℃以下に冷却して焼鈍した後、圧下率６０〜９０％の最終冷間圧延を施し、或いは熱間仕上圧延後焼鈍しないで圧下率６０〜９０％の最終冷間圧延を施し、その後必要に応じ１００〜１５０℃の温度で仕上焼鈍を施すことを特徴とする耳率の低いキャンボディ用アルミニウム合金板の製造方法である。
【０００７】
【発明の実施の形態】
以下に本発明にて製造するアルミニウム合金板の合金組成について説明する。
Ｍｇは強度向上に寄与し、特に缶底部の高強度化に有効である。その含有量を０．８〜１．４wt％に限定する理由は、０．８wt％未満ではその効果が十分に得られず、１．４wt％を超えるとＤＩ成形時に加工硬化し易くなり、しごき加工時の割れの発生頻度が増加するためである。Ｍｇの最適含有量は、他元素の添加量や製造条件によりやや変化するが、強度とＤＩ(Drawn and Ironing) 成形性のバランスが良好な組成範囲は１．０〜１．３５wt％で、さらに望ましくは１．1 〜１．３wt％の範囲である。
【０００８】
Ｍｎは強度とＤＩ成形性の向上に寄与する。
ＭｎがＤＩ成形性を向上させるのは、Ｍｎが固体潤滑作用を有するＡｌ−Ｍｎ系、Ａｌ−Ｍｎ−Ｆｅ系、Ａｌ−Ｍｎ−Ｆｅ−Ｓｉ系等の晶出物を形成するためである。すなわち、ＤＩ成形には、通常エマルジョン型の潤滑剤が使用されるが、これだけでは潤滑が不十分であり、アルミニウム合金板と金型との凝着によるビルトアップが発生してゴーリング又はスコアリングと呼ばれる擦り傷や焼付きが発生することがある。Ｍｎは前記晶出物を形成することでビルトアップの発生を抑制する。Ｍｎの含有量を０．７〜１．３wt％に限定する理由は、０．７wt％未満ではＤＩ成形性の改善効果が不十分なばかりか強度も不足し、１．３wt％を超えるとＤＩ成形性および強度向上効果が飽和する上、溶解鋳造時に、後述するＦｅと反応してＡｌ−Ｍｎ−Ｆｅ系の巨大な（時として数ｍｍ程度のサイズの）初晶化合物を形成してＤＩ成形時に割れやピンホールを誘発するためである。Ｍｎの望ましい含有量は０．９〜１．２wt％、さらに望ましくは１．０〜１．２wt％である。
【０００９】
Ｆｅは前記Ｍｎの晶出物の生成を促進するとともにその分布状態を均一化してＤＩ成形性をより一層向上させる。Ｆｅの含有量を０．２〜０．５wt％に限定する理由は、０．２wt％未満ではその効果が十分に得られず、０．５wt％を超えると前述のＡｌ−Ｍｎ−Ｆｅ系の巨大初晶化合物が発生し易くなるためである。Ｆｅの望ましい含有量は０．３〜０．５wt％、さらに望ましくは０．３５〜０．４５wt％である。
【００１０】
Ｓｉは、Ａｌ−Ｍｎ−Ｆｅ系の晶出物に相変態を起こさせ、より硬度の高いＡｌ−Ｍｎ−Ｆｅ−Ｓｉ系析出物を形成して、しごき加工性を向上させる。Ｓｉの含有量を０．１〜０．５wt％に限定する理由は、０．１wt％未満ではその効果が十分に得られず、０．５wt％を超えると晶出物が巨大化して、逆にしごき加工性が低下するためである。
【００１１】
ＣｕはＭｇと同じように缶底部の高強度化に有効である。
Ｃｕの含有量を０．１〜０．３wt％に限定する理由は、０．１wt％未満ではその効果が十分に得られず、また耐圧強度を確保するために必要な最終冷間圧延での圧下率が大きくなってＤＩ成形性が低下し、０．３wt％を超えると合金板は加工硬化し易くなり、しごき加工時に硬化して、逆にしごき加工性が低下するためである。
【００１２】
Ｔｉ、又はＴｉおよびＢは、鋳塊の結晶粒を均一に微細化して成形加工性を改善する。
Ｔｉの含有量を０．００５〜０．０５wt％に限定する理由は、０．００５wt％未満ではその効果が十分に得られず、０．０５wt％を超えるとＡｌ−Ｔｉ系の巨大双晶化合物が溶解鋳造時に発生し易くなり、これが圧延後も残存してＤＩ成形時に割れやピンホールの発生原因になるためである。
【００１３】
ＢはＴｉの結晶粒微細化効果を助長する。Ｂが０．０００１wt％未満ではその効果が十分に得られず、０．０１wt％を超えるとＴｉ−Ｂ系の巨大な双晶化合物が溶解鋳造時に発生し易くなり、これが圧延後も残存して成形時における割れやピンホールの発生頻度が増加する原因になる。
不純物については、本発明を損なわない程度で許容できる。例えばＺｎは０．３wt％以下、Ｃｒは０．３wt％以下、Ｚｒは０．１wt％以下、Ｖは０．１wt％以下であれば問題ない。
【００１４】
次に本発明の製造方法について説明する。
前記組成のアルミニウム合金を、例えば、通常のＤＣ鋳造法（半連続鋳造法）により鋳造し、得られた鋳塊を所定温度で均質化処理する。この均質化処理条件を５６０〜６２０℃で１時間以上に限定する理由は、均質化処理温度が５６０℃未満でも１時間未満でも十分に均質化されず、６２０℃を超えると鋳塊表面に膨れが生じるためである。生産性とその効果を勘案した最も望ましい均質化処理条件は５６０〜６２０℃で３〜１２時間加熱する条件である。
【００１５】
本発明では、均質化処理後、放冷し鋳塊表面に存在する偏析層、酸化膜等を切削除去した後、再び適当な温度に加熱して熱間粗圧延と仕上圧延を施す。
この熱間粗圧延と熱間仕上圧延工程で歪み（再結晶駆動力）を多く蓄積させることにより、マトリクス中の遷移帯(transition band) から核生成し成長する１種の再結晶集合組織である立方体方位が優先的に生じた組織を形成させる。
熱間粗圧延を４５０〜５５０℃に加熱して行う理由は、４５０℃未満では十分な圧延加工性が得られず、５５０℃を超えると粗圧延板の表面が酸化したり、再結晶粒が粗大化して成形性が低下するためである。
本発明では均質化処理後、２０℃／時間以上の冷却速度で４５０〜５５０℃まで冷却し、続いて熱間粗圧延を施すか、或いはそのまま室温まで冷却後３０℃／時間以上の昇温速度で４５０〜５５０℃に再加熱して熱間粗圧延を施す。
【００１６】
前記均質化処理後熱間粗圧延温度までの冷却速度を２０℃／時間以上にするのは、表面酸化や結晶粒が成長し易い５５０℃を超える温度域を素早く通過させて表面酸化などを抑えるためである。
また室温から熱間粗圧延までの昇温速度を３０℃／時間以上にするのは、析出物の個数密度（単位体積当たりの析出物の個数）が急激に増加する４５０℃未満の温度域を素早く通過させて前記析出物の増加を抑えるためである。前記均質化処理温度から室温までの間、および室温から４５０℃に達するまでの間に生成する析出物は微細で耳率に悪影響を及ぼす。すなわちこの微細析出物は後の焼鈍過程において、０°−９０°耳成分となる立方体方位再結晶粒の成長を妨げる。
【００１７】
本発明において、熱間粗圧延を終了板厚１２〜５０ｍｍに限定する理由は、熱間粗圧延終了板厚が１２ｍｍ未満では熱間仕上圧延に入る前に熱延板が冷えてしまい、所望の粗圧延終了温度（３００〜４５０℃) が得難くなり、５０ｍｍを超えると熱延板の表面性状（焼付き、肌荒れ等）を悪化させずに熱間仕上圧延での最終板厚を１．６〜３．０ｍｍにすることが難しくなるためである。
また熱間粗圧延終了温度は３００〜４５０℃に規定する理由は、熱間粗圧延終了温度が３００℃未満では熱間仕上圧延開始温度が低くなりすぎて仕上圧延時にエッジ部が割れる等の問題が生じ、前記終了温度が４５０℃を超えると粗圧延終了時の再結晶率が３０％を超える程度に多くなり、後の工程で仕上圧延を行っても歪を十分蓄積できず、仕上圧延後室温まで冷却した後の再結晶で立方体方位を優先的に生じた組織にすることができないためである。特に望ましい粗圧延終了温度は３３０〜３８０℃である。
また熱間粗圧延の最終パス圧下率Ｒを〔７０−０．２Ｓ（Ｓ：圧延速度ｍ／分) 〕％以下にする理由は、前記圧下率Ｒが前記圧下率を超えると、圧延材が加工発熱して４５０℃を超えるとともに、粗圧延での歪み（再結晶駆動力）が大きくなって、熱間粗圧延終了後に再結晶率が３０％を超えてしまうためである。
また熱間粗圧延終了後熱間仕上圧延開始までの時間ｔ秒〔ｔ＝２．８×１０⁴exp（−０．０１２Ｔ），Ｔ：熱間粗圧延終了温度℃) 〕以内に行う理由は、ｔ秒を超えると歪みが回復し、再結晶率が３０％を超えてしまうためである。
【００１８】
本発明において、熱間仕上圧延を、スタンド数３以上のタンデム式熱間仕上圧延機を用い、総圧下率を８０％以上とし、各スタンドでの圧下率を３０％以上の条件で行う理由は、前記仕上圧延機のスタンド数が３未満でも、総圧下率が８０％未満でも、各スタンドでの圧下率が３０％未満でも、歪みの蓄積が不十分であり、熱間仕上圧延後に立方体方位再結晶粒を得るための駆動力が不足し耳率が増加するためである。
また熱間仕上圧延の終了板厚１．６〜３．０ｍｍとする理由は、前記終了板厚が１．６ｍｍ未満では、熱延板の表面性状（焼付き、肌荒れなど）および板厚分布が悪化し、３．０ｍｍを超えると後工程の最終冷間圧延で圧下率が高くなって、耳率の低いアルミニウム合金板を得ることが困難になるためである。
また熱間仕上圧延の終了温度を２９０℃以上に限定する理由は、前記終了温度が２９０℃未満では、熱間仕上圧延終了後の再結晶率が８０％未満となって立方体方位優先の再結晶集合組織が十分発達しないためである。
なお、その後に仕上焼鈍を施して再結晶率を８０％以上に高めても０°−９０°耳を低下させる立方体方位以外の方位（例えばＲ方位）も発達するため効果がない。この傾向は焼鈍を施さずに最終冷間圧延を行う場合に一層強く現れる。
【００１９】
本発明では、熱間仕上圧延後、そのまま、または中間焼鈍（箱型焼鈍、連続焼鈍）後に、圧下率６０〜９０％の最終冷間圧延を施す。
この最終冷間圧延により缶胴材として必要な缶強度が付与される。前記冷間圧延の圧下率を６０〜９０％に限定する理由は、６０％未満では十分な耐圧強度が得られず、９０％を超えると深絞り成形時の４５°耳の耳率が高くなるとともに、強度が高くなりすぎてＤＩ成形性が低下し、カッピング割れ、缶底割れの発生頻度が高くなるためである。この最終冷間圧延の終了板厚は０．２８〜０．４ｍｍである。
【００２０】
本発明において、箱型焼鈍炉を用いて行う焼鈍を、３００〜４５０℃の温度に３０分以上保持して行う理由は、前記焼鈍温度が３００℃未満でも、焼鈍時間が３０分未満でも、完全再結晶組織が十分得られず、４５０℃を超えると再結晶した結晶粒が粗大に成長し、この粗大再結晶組織は加工性を低下させる危険があると同時に、特定方位の結晶粒が優先的に成長して冷間圧延板の４５°耳を大きくする場合があるためである。
【００２１】
本発明において、連続焼鈍炉を用いて行う焼鈍を、１００℃／分以上の加熱速度で３６０〜５６０℃の温度に加熱し、前記温度に到達後直ちに或いは１２０秒以下の時間保持後１００℃／分以上の冷却速度で７０℃以下に冷却して行う理由は、前記焼鈍温度が３６０℃未満では、再結晶が不十分なため冷間圧延板の強度が上がりすぎてＤＩ成形性が低下し、５６０℃を超えるとＣｕやＳｉ等の析出物が再固溶しすぎて、これが塗装焼付け時に析出してフランジ成形性が低下し、また保持時間が１２０秒を超えると、焼鈍温度が５６０℃以下でも析出物が再固溶しすぎ、この再固溶元素（ＣｕやＳｉなど）が塗装焼付け時に析出してフランジ成形性を低下させる。保持時間は０でも良い。すなわち目標温度に到達後直ちに冷却してもよい。
加熱および冷却速度をともに１００℃／分以上にしたのは生産性を高めるためである。冷却速度の場合は、１００℃／分未満では、固溶したＣｕおよびＳｉが析出して次の最終冷間圧延で十分な強度が得られなくなるためでもある。
【００２２】
熱間圧延終了後直ちに、或いは焼鈍後に最終冷間圧延を施す理由は、この最終冷間圧延により、缶胴材として必要な強度を付与するためである。この最終冷間圧延の終了板厚は通常０．２８〜０．４ｍｍである。この最終冷間圧延での圧下率を６０〜９０％に限定する理由は、６０％未満では合金板の強度が低く耐圧強度が不足し、９０％を超えると、深絞り成形時の４５°耳の耳率が高くなるとともに、冷間圧延板強度が高くなりすぎてＤＩ成形性が低下し、カッピング割れ、缶底割れの発生頻度が高くなるためである。
【００２３】
最終冷間圧延後、必要に応じて、仕上焼鈍を施す。
この仕上焼鈍により加工組織に回復が起きて、カッピング成形性や缶底成形性が向上する。前記仕上焼鈍を１００〜１５０℃の温度で行う理由は、１００℃未満ではその効果が十分に得られず、１５０℃を超えると、固溶元素が析出しすぎてＤＩ成形性やフランジ成形性が低下するようになる。最も望ましい仕上焼鈍条件は１１５〜１５０℃である。前記仕上焼鈍時間は８時間以下、特に望ましくは１〜４時間である。
【００２４】
【実施例】
以下に本発明を実施例により詳細に説明する。
（実施例１）
表１に示す本発明組成のＡｌ合金組成Ａ〜Ｅ、および比較例として本発明と組成の異なるＡｌ合金Ｆ〜Ｍを常法により溶解鋳造して、厚さ５００ｍｍの鋳塊（スラブ）を得た。次にこの鋳塊を厚さ４９０ｍｍに面削し、次いで６００℃で６時間の均質化処理を施し、その後３５０℃まで放冷し、そこから室温まで水冷し、次いで昇温速度５０℃／時間で５２０℃まで再加熱して熱間粗圧延を行った。前記熱間粗圧延は終了板厚２５ｍｍ、最終パスの圧下率２５％、圧延速度１２０ｍ／分、圧延終了温度３６０℃の条件で行った。
熱間粗圧延が終了した１８０秒後に熱間仕上圧延を開始し、厚さ２．２ｍｍの熱延板を得た。熱間仕上圧延は４スタンドの仕上圧延機を用い、総圧下率９０．４％、各スタンドでの圧下率Ｆ₁:５３％、Ｆ₂:４７％、Ｆ₃:４４％、Ｆ₄:３７％、圧延終了温度３３０℃の条件で行った。再結晶率は熱間仕上圧延前が５％、圧延終了後が１００％であった。前記再結晶率とは熱延板断面に占める再結晶粒の面積比率である。
熱間仕上圧延終了後、連続焼鈍炉により４００℃で０分（材料４００℃に到達後直ちに空冷）焼鈍した。このときの加熱速度は８５０℃／分、冷却速度は１０００℃／分とした。続いて常法により板厚０．３ｍｍまで最終冷間圧延 (最終冷間圧下率８７．５％）し、次いで１１５℃で２時間の仕上焼鈍を施して缶胴用Ａｌ合金板を製造した。
【００２５】
このようにして得られた各々のＡｌ合金板について、耳率、引張強度（引張強さ（ＴＳ）と0.2%耐力（ＹＳ））、ＤＩ成形性、フランジ成形性を調査した。
耳率は、直径３３ｍｍ、肩Ｒ２．５ｍｍのポンチを用いて５７ｍｍφの円板をクリアランス３０％で深絞りして測定した。
引張強度は、２００℃で２０分間加熱（塗装焼付け条件）後にも測定した。
ＤＩ成形性は、炭酸飲料用のＤＩ缶胴（内径６６ｍｍφ、側壁板厚１０３μｍ、側壁先端部板厚１６５μｍ) に成形して調査した。
フランジ成形性は、前記成形したＤＩ缶をトリミングと洗浄後、２００℃で２０分間加熱し、次いで４段ネッキング加工を施して開口部の内径ｄを５７ｍｍφに縮小し、最後に頂角９０°の円錐状の治具の頂部を割れが発生するまで押し込み、割れが発生した時の開口部の径Ｄを測定し、開口部の径の増加率Ｐを、Ｐ＝〔（Ｄ−ｄ）／ｄ〕×１００％の式により計算して評価した。
結果を表２に示す。
【００２６】
【表１】

【００２７】
【表２】

【００２８】
表２より明らかなように、本発明例品の No.Ａ〜Ｅは耳率が２．５％以下と低く、フランジ成形での口径の限界増加率が１．５％以上と大きくフランジ成形性が良好であった。また２００℃で２０分間加熱後の耐力（ＹＳ）も２５０ＭＰａ以上あり、缶底部の耐圧強度も問題のない水準であった。またＤＩ成形性も良好であった。
これに対し、比較例品の No.Ｆと No.ＧはそれぞれＭｇまたはＭｎの含有量が多かったために２００℃で２０分間の加熱により引張強さが高くなり、このため缶胴側壁先端部の塗装焼付け加熱による軟化が不十分となりＤＩ成形でしごき割れが生じた。No. ＨはＳｉ、Ｃｕが多いためフランジ成形性が劣った。 No.ＩはＭｇの含有量が少ないため強度が低下し、 No.ＪはＭｎの含有量が少ないためＤＩ成形において焼付きが生じた。 No.ＫはＣｕとＳｉの含有量が少ないため強度が低下した。No. Ｌ、Ｍは不純物のＺｎまたはＣｒが多いため強度が高くなり、ＤＩ成形でしごき割れが生じた。
【００２９】
（実施例２）
表１に示した No.ＡのＡｌ合金を常法により溶解鋳造して厚さ５００ｍｍの鋳塊（スラブ）を得た。次にこの鋳塊を厚さ４９０ｍｍまで面削し、次いで均質化処理、冷却、加熱処理、熱間粗圧延、熱間仕上圧延を順に施して熱延コイルを得た。この熱延コイルを室温まで冷却した後、箱型焼鈍または連続焼鈍により焼鈍し、または焼鈍を行わずに、引続き常法により冷間圧延して缶胴用Ａｌ合金板を製造した。均質化処理、熱延、焼鈍、最終冷間圧延の条件は表３、５に示すように種々に変化させた。なお、本発明と異なる条件で実施した比較例を表４、６に示す。
【００３０】
【表３】

【００３１】
【表４】

【００３２】
【表５】

【００３３】
【表６】

【００３４】
このようにして得られた各々のＡｌ合金板について、実施例１と同じ方法により、耳率、引張強度、ＤＩ成形性、フランジ成形性を調査した。耳率（％）は、直径３３ｍｍ、肩Ｒ２．５ｍｍのポンチを用いて５７ｍｍφの円板をクリアランス３０％で深絞りして測定した。表３〜６に製造条件、表７、８に測定結果を示す。評価基準は、耳率２．５％以内、加熱処理（２００℃×２０分) 後の耐力２６０ＭＰａ以上、フランジ成形での口径の限界増加率１５％以上を良好とした。
【００３５】
【表７】

【００３６】
【表８】

【００３７】
表７より明らかなように本発明例 (No.1〜11) は耳率が２．５％以下と低く、フランジ成形性も良好であった。また塗装焼付けに相当する加熱処理後の強度（耐力）も２５０ＭＰａ以上で、缶底部の耐圧性にも問題のない強度水準を有し、更にＤＩ成形性も良好であった。
【００３８】
これに対し、比較例はいずれも、表８より明らかなように、何らかの特性が不良であった。
すなわち、比較例のNo.1は均質化処理温度が低かったため、均質化が不十分で耳率が高くなった。
No.2は均質化処理温度が高く鋳塊表面に膨れが生じ、仕上圧延終了後の表面性状が悪化した。
No.3,4は冷却速度または昇温速度が遅かったために、析出物の個数密度が増加し耳率が基準値を上回った。
No.5は昇温温度が低かったために、析出物の個数密度が増加し耳率が基準値を上回った。
No.6,7は熱間粗圧延最終パスＲ、熱間粗圧延終了後熱間仕上圧延開始までの時間ｔ秒が本発明の規定値から外れており、粗圧延終了から仕上圧延開始までの再結晶率が３０％を超えたため歪みの蓄積が不十分で耳率が基準値を上回った。
No.8は終了板厚が厚かったため、最終冷間圧下率が高くなり、その結果ＤＩ成形で絞り割れが発生し、耳率が基準値を上回った。
No.9は終了板厚が薄く熱間仕上圧延後に焼付きが生じ、缶に成形したときの缶表面にキズが発生した。
No.10 は熱間仕上圧延終了温度が低すぎたために、終了後の再結晶率が低く、耳率が基準値を上回った。
No.11,12は熱間仕上圧延での各パス圧下率が３０％未満であり、熱間仕上げ圧延での総圧下率が８０％未満であるために、歪みの蓄積が不十分で耳率が基準値を上回った。
No.13 は最終冷間圧下率が高くなったためにＤＩ成形で絞り割れが発生し、耳率が基準値を上回った。
No.14 は仕上焼鈍温度が高いために析出により強度（焼付け前の耐力）が高くなり、しごき割れが発生した。
No.15 は中間焼鈍温度が本発明の条件より高かったため２００℃で２０分の焼付けによる引張強さが向上し、従って缶胴側壁先端部分の塗装焼付け加熱による熱軟化が起こらないので、フランジ成形での口径の限界増加率が小さくフランジ成形性が劣った。
【００３９】
【発明の効果】
以上に述べたように、本発明によれば、強度、しごき加工性、塗装焼付け後のフランジ成形性に優れた耳率の低いキャンボディ用Ａｌ合金板が得られ、工業上顕著な効果を奏する。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for producing an aluminum alloy plate for a can body, which is particularly suitable for beverage can body materials, has high strength, excellent ironing workability, and excellent flange formability after painting and baking, and has a low ear rate.
[0002]
[Prior art]
Conventionally, beverage can bodies are manufactured by homogenizing a JIS-3004 alloy ingot and then hot rolling, cold rolling, annealing, and cold rolling. After cold rolling, finish annealing, degreasing, washing, application of lubricating oil for cupping, etc. are performed as necessary.
[0003]
When the beverage can body is squeezed into a cup (cylindrical) shape, irregularities may occur on the peripheral edge of the cup. The ratio of the convex and concave height to the cup height is referred to as an ear ratio, but when the ear ratio is high, a chip jumps from the tip of the ear during cup molding and iron molding, and a pinhole or tear-off occurs, The dimensional accuracy of the can after flange molding may be reduced. Therefore, when the ear ratio is high, the amount of trimming after molding the can body is increased, but there is a problem that the recess remains after trimming. In addition, for cans with small diameters that have recently increased, if a material with a high ear rate is used, there will be a large variation in the length of the flange during flanging and necking, which will hinder the tightening of the lid. Is causing problems. For this reason, a beverage can body material is required to have a material with a lower ear ratio than ever before.
[0004]
By the way, the ears are caused by the crystallographic anisotropy of the rolled material, and the height of the ears is the reorientation of cube-oriented recrystallized grains formed by recrystallization that proceeds after the end of hot rolling or during annealing. It is determined by the balance between the texture component (mainly 0 ° -90 ° ears) and the rolling texture component (45 ° ears) formed by rolling (cold rolling).
For example, when emphasizing the strength of the can, since cold working is performed at a high pressure ratio, a rolling texture is strongly formed. Therefore, in this case, the hot rolling or annealing conditions are strictly defined and the cubic orientation recrystallized grains are preferentially grown (JP-A-4-228551, JP-A-6-158244).
However, in recent years, as the can diameter has decreased, the user's requirements for the ear rate have become increasingly severe, and the low ear rate required by the user cannot be realized simply by specifying the hot rolling conditions or annealing conditions. In order to further lower the temperature, in addition to hot rolling conditions, it has been necessary to extensively study the homogenization conditions of the upper process.
[0005]
[Problems to be solved by the invention]
An object of the present invention is to provide a method for producing an aluminum alloy plate for a can body that can sufficiently and reliably reduce the ear rate even for a reduced diameter tube and is excellent in formability and strength.
[0006]
[Means for Solving the Problems]
In the present invention, Mg is 0.8 to 1.4 wt%, Mn is 0.7 to 1.3 wt%, Fe is 0.2 to 0.5 wt%, Si is 0.1 to 0.5 wt%, Cu is An aluminum alloy ingot containing 0.1 to 0.3 wt%, Ti 0.005 to 0.05 wt% alone or together with B0.0001 to 0.01 wt%, with the balance being Al and inevitable impurities, Apply homogenization treatment for 1 hour or more in a temperature range of 620 ° C., then cool to 450 to 550 ° C. at a cooling rate of 20 ° C./hour or more from the homogenization treatment temperature, and perform hot rough rolling or homogenization After the treatment, it is cooled as it is to room temperature, and then reheated to 450 to 550 ° C. at a temperature rising rate of 30 ° C./hour or more to perform hot rough rolling. The hot rough rolling is finished with a plate thickness of 12 to 50 mm and an end temperature. Is 300 to 450 ° C., final pass reduction ratio R is R ≦ 70−0.2S S: subjected under the conditions of the rolling speed m / min), rough hot rolling finished after t seconds ^{(t = 2.8 × 10 4 exp} (-0.012T), T: hot rough rolling finishing temperature ° C.) within Hot finish rolling was started, and the hot finish rolling was performed using a tandem hot finish rolling mill having three or more stands. The total reduction rate was 80% or more, the reduction rate at each stand was 30% or more, and the finished plate thickness was 1 .6 ~ 3.0mm, finish temperature is 290 ℃ or more, and after hot finish rolling, cool to room temperature, then hold at 300-450 ℃ for 30 minutes or more using a box-type annealing furnace for annealing In addition, using a continuous annealing furnace, annealing is performed at a temperature rising rate of 100 ° C./min or more at a temperature of 360 to 560 ° C., and immediately after reaching the temperature or after holding for a time of 120 seconds or less, 100 ° C./min or more. After cooling to 70 ° C. or less at the cooling rate and annealing, the final cold rolling with a rolling reduction of 60 to 90% Alternatively, the final cold rolling with a reduction ratio of 60 to 90% is performed without annealing after hot finish rolling, and then finish annealing is performed at a temperature of 100 to 150 ° C. as necessary. It is a manufacturing method of the aluminum alloy plate for bodies.
[0007]
DETAILED DESCRIPTION OF THE INVENTION
Below, the alloy composition of the aluminum alloy plate manufactured by this invention is demonstrated.
Mg contributes to strength improvement and is particularly effective for increasing the strength of the bottom of the can. The reason for limiting the content to 0.8 to 1.4 wt% is that if less than 0.8 wt%, the effect cannot be sufficiently obtained, and if it exceeds 1.4 wt%, it becomes easy to work and harden during DI molding, and ironing This is because the frequency of occurrence of cracks during processing increases. The optimum content of Mg varies slightly depending on the addition amount of other elements and production conditions, but the composition range with a good balance between strength and DI (Drawn and Ironing) moldability is 1.0 to 1.35 wt%, Desirably, it is in the range of 1.1 to 1.3 wt%.
[0008]
Mn contributes to improvement of strength and DI moldability.
Mn improves DI moldability because Mn forms crystallized substances such as Al-Mn, Al-Mn-Fe, and Al-Mn-Fe-Si that have a solid lubricating action. That is, an emulsion type lubricant is usually used for DI molding, but this alone is insufficient in lubrication, and build-up due to adhesion between the aluminum alloy plate and the mold occurs, resulting in goling or scoring. Scratches and seizures sometimes occur. Mn suppresses the occurrence of built-up by forming the crystallized product. The reason for limiting the Mn content to 0.7 to 1.3 wt% is that if it is less than 0.7 wt%, not only the effect of improving the DI moldability is insufficient but also the strength is insufficient. In addition to saturation of moldability and strength improvement effect, it reacts with Fe, which will be described later, at the time of melt casting to form an Al-Mn-Fe-based huge (sometimes several mm in size) primary compound and DI molding This is to sometimes induce cracks and pinholes. The desirable content of Mn is 0.9 to 1.2 wt%, more preferably 1.0 to 1.2 wt%.
[0009]
Fe promotes the formation of Mn crystallized substances and makes the distribution state uniform to further improve the DI moldability. The reason for limiting the Fe content to 0.2 to 0.5 wt% is that the effect cannot be sufficiently obtained if the content is less than 0.2 wt%, and if the content exceeds 0.5 wt%, the above-described Al-Mn-Fe system is used. This is because a large primary crystal compound is likely to be generated. A desirable content of Fe is 0.3 to 0.5 wt%, and more desirably 0.35 to 0.45 wt%.
[0010]
Si causes a phase transformation in the Al-Mn-Fe-based crystallized substance, and forms an Al-Mn-Fe-Si-based precipitate having a higher hardness, thereby improving ironing workability. The reason for limiting the Si content to 0.1 to 0.5 wt% is that the effect cannot be sufficiently obtained if it is less than 0.1 wt%, and if the content exceeds 0.5 wt%, the crystallized material becomes enormous, This is because the ironing processability is lowered.
[0011]
Cu is effective in increasing the strength of the bottom of the can as in the case of Mg.
The reason for limiting the Cu content to 0.1 to 0.3 wt% is that the effect is not sufficiently obtained if it is less than 0.1 wt%, and in the final cold rolling necessary to ensure the compressive strength. This is because when the rolling reduction increases, the DI formability decreases, and when it exceeds 0.3 wt%, the alloy plate is easy to work harden and harden during the ironing process, and conversely the iron workability decreases.
[0012]
Ti, or Ti and B improve the moldability by uniformly refining the crystal grains of the ingot.
The reason for limiting the Ti content to 0.005 to 0.05 wt% is that the effect cannot be sufficiently obtained if it is less than 0.005 wt%, and if it exceeds 0.05 wt%, an Al—Ti giant twin compound This is because it tends to occur during melt casting, which remains after rolling and causes cracks and pinholes during DI molding.
[0013]
B promotes the grain refinement effect of Ti. If B is less than 0.0001 wt%, the effect cannot be sufficiently obtained. If it exceeds 0.01 wt%, a Ti-B giant twin compound is likely to be generated during melt casting, and this remains after rolling. This can increase the frequency of cracks and pinholes during molding.
Impurities are acceptable to the extent that the present invention is not impaired. For example, there is no problem if Zn is 0.3 wt% or less, Cr is 0.3 wt% or less, Zr is 0.1 wt% or less, and V is 0.1 wt% or less.
[0014]
Next, the manufacturing method of this invention is demonstrated.
The aluminum alloy having the above composition is cast, for example, by a normal DC casting method (semi-continuous casting method), and the resulting ingot is homogenized at a predetermined temperature. The reason for limiting the homogenization treatment condition to 560 to 620 ° C. for 1 hour or longer is that the homogenization treatment temperature is less than 560 ° C. or less than 1 hour, and is not sufficiently homogenized. This is because. The most desirable homogenization treatment conditions in consideration of productivity and its effect are the conditions of heating at 560-620 ° C. for 3-12 hours.
[0015]
In the present invention, after the homogenization treatment, the mixture is allowed to cool and the segregation layer, oxide film and the like existing on the ingot surface are cut and removed, and then heated to an appropriate temperature and subjected to hot rough rolling and finish rolling.
It is a recrystallization texture that nucleates and grows from the transition band in the matrix by accumulating a lot of strain (recrystallization driving force) in this hot rough rolling and hot finish rolling process. A structure in which cube orientation is preferentially formed is formed.
The reason why hot rough rolling is performed by heating to 450 to 550 ° C. is that sufficient rolling workability is not obtained at less than 450 ° C., and when the temperature exceeds 550 ° C., the surface of the rough rolled plate is oxidized or recrystallized grains are not formed. This is because it becomes coarse and the moldability deteriorates.
In the present invention, after the homogenization treatment, it is cooled to 450 to 550 ° C. at a cooling rate of 20 ° C./hour or more, followed by hot rough rolling, or after being cooled to room temperature as it is, a heating rate of 30 ° C./hour or more. To 450-550 ° C. and hot rough rolling is performed.
[0016]
The reason why the cooling rate to the hot rough rolling temperature after the homogenization treatment is set to 20 ° C./hour or more is to allow surface oxidation and crystal grains to grow easily to quickly pass through a temperature range exceeding 550 ° C. to suppress surface oxidation and the like. Because.
The temperature increase rate from room temperature to hot rough rolling is 30 ° C./hour or more because the number density of precipitates (the number of precipitates per unit volume) increases rapidly in a temperature range below 450 ° C. It is for passing through quickly and suppressing the increase of the said deposit. Precipitates formed between the homogenization temperature and room temperature and between room temperature and 450 ° C. are fine and adversely affect the ear rate. That is, this fine precipitate hinders the growth of cubic orientation recrystallized grains that become 0 ° -90 ° ear components in the subsequent annealing process.
[0017]
In the present invention, the reason for limiting the hot rough rolling to the finished sheet thickness of 12 to 50 mm is that if the hot rough rolled finished sheet thickness is less than 12 mm, the hot rolled sheet is cooled before entering the hot finish rolling. The rough rolling finish temperature (300 to 450 ° C.) becomes difficult to obtain, and if it exceeds 50 mm, the surface thickness (seizure, rough surface, etc.) of the hot rolled sheet is not deteriorated and the final thickness in hot finish rolling is 1.6. It is because it becomes difficult to make it to -3.0 mm.
Moreover, the reason for prescribing the hot rough rolling end temperature to 300 to 450 ° C. is that if the hot rough rolling end temperature is less than 300 ° C., the hot finish rolling start temperature becomes too low and the edge portion breaks during finish rolling. When the end temperature exceeds 450 ° C., the recrystallization rate at the end of rough rolling increases to a level exceeding 30%, and strain cannot be sufficiently accumulated even after finish rolling in the subsequent process. This is because it is impossible to obtain a structure in which the cubic orientation is preferentially generated by recrystallization after cooling to room temperature. A particularly desirable rough rolling end temperature is 330 to 380 ° C.
The reason for setting the final pass rolling reduction R of hot rough rolling to [70-0.2S (S: rolling speed m / min)]% or less is that when the rolling reduction R exceeds the rolling reduction, the rolled material This is because the heat generated by the process exceeds 450 ° C., and the distortion (recrystallization driving force) in the rough rolling increases, and the recrystallization rate exceeds 30% after the hot rough rolling.
Further, the reason for performing within t seconds [t = 2.8 × 10 ⁴ exp (−0.012 T), T: hot rough rolling finish temperature ° C.)] from the end of hot rough rolling to the start of hot finish rolling is as follows. This is because the strain recovers after t seconds and the recrystallization rate exceeds 30%.
[0018]
In the present invention, the hot finish rolling is performed using a tandem hot finish rolling mill having three or more stands, the total rolling reduction is 80% or more, and the rolling reduction at each stand is 30% or more. Even if the number of stands of the finish rolling mill is less than 3, the total rolling reduction is less than 80%, or the rolling reduction at each stand is less than 30%, the accumulation of strain is insufficient, and the cube orientation after hot finish rolling This is because the driving force for obtaining recrystallized grains is insufficient and the ear rate increases.
Moreover, the reason why the finish plate thickness of hot finish rolling is 1.6 to 3.0 mm is that when the finish plate thickness is less than 1.6 mm, the surface properties (seizure, rough skin, etc.) and the plate thickness distribution of the hot rolled plate are This is because if the thickness is more than 3.0 mm, the reduction ratio becomes high in the final cold rolling in the subsequent step, and it becomes difficult to obtain an aluminum alloy sheet having a low ear ratio.
Moreover, the reason for limiting the finish temperature of hot finish rolling to 290 ° C. or higher is that if the finish temperature is less than 290 ° C., the recrystallization rate after the finish of hot finish rolling is less than 80%, giving priority to cube orientation recrystallization. This is because the texture does not develop sufficiently.
Note that even if finish annealing is performed thereafter to increase the recrystallization rate to 80% or more, there is no effect because orientations other than the cubic orientation (for example, R orientation) that lower the 0 ° -90 ° ears develop. This tendency appears even more strongly when the final cold rolling is performed without annealing.
[0019]
In the present invention, after the hot finish rolling, as it is or after intermediate annealing (box annealing, continuous annealing), final cold rolling with a rolling reduction of 60 to 90% is performed.
This final cold rolling provides the can strength necessary for the can body material. The reason for limiting the rolling reduction of the cold rolling to 60 to 90% is that sufficient pressure strength is not obtained if it is less than 60%, and the ear rate of 45 ° ears during deep drawing is increased if it exceeds 90%. At the same time, the strength becomes too high, the DI moldability is lowered, and the occurrence frequency of cupping cracks and can bottom cracks is increased. The final thickness of the final cold rolling is 0.28 to 0.4 mm.
[0020]
In the present invention, the reason why the annealing performed using a box-type annealing furnace is maintained at a temperature of 300 to 450 ° C. for 30 minutes or more is that the annealing temperature is less than 300 ° C. and the annealing time is less than 30 minutes. A recrystallized structure cannot be obtained sufficiently, and when it exceeds 450 ° C, the recrystallized crystal grains grow coarsely. This coarse recrystallized structure has a risk of lowering workability, and at the same time, crystal grains with a specific orientation are preferential. This is because the 45 ° ear of the cold-rolled sheet may be enlarged.
[0021]
In the present invention, annealing performed using a continuous annealing furnace is heated to a temperature of 360 to 560 ° C. at a heating rate of 100 ° C./min and immediately after reaching the temperature or after holding for a time of 120 seconds or less, 100 ° C. / The reason for cooling to 70 ° C. or less at a cooling rate of at least minutes is that when the annealing temperature is less than 360 ° C., the recrystallization is insufficient and the strength of the cold-rolled sheet is too high, and the DI moldability is lowered. When it exceeds 560 ° C, precipitates such as Cu and Si are re-dissolved too much, and this precipitates during coating baking, and flange formability deteriorates. When the holding time exceeds 120 seconds, the annealing temperature is 560 ° C or less. However, the precipitate is re-dissolved too much, and this re-dissolved element (Cu, Si, etc.) is precipitated during coating baking to reduce the flange formability. The holding time may be zero. That is, it may be cooled immediately after reaching the target temperature.
The reason why both the heating and cooling rates are set to 100 ° C./min or more is to increase productivity. In the case of the cooling rate, if it is less than 100 ° C./min, solid solution of Cu and Si is precipitated, and sufficient strength cannot be obtained in the next final cold rolling.
[0022]
The reason why the final cold rolling is performed immediately after the end of hot rolling or after annealing is to give the necessary strength as a can body material by this final cold rolling. The final plate thickness of the final cold rolling is usually 0.28 to 0.4 mm. The reason for limiting the rolling reduction in this final cold rolling to 60 to 90% is that if it is less than 60%, the strength of the alloy sheet is low and the pressure resistance is insufficient, and if it exceeds 90%, it is 45 ° during deep drawing. This is because the ear rate is increased, the cold rolled sheet strength is excessively increased, the DI moldability is decreased, and the occurrence frequency of cupping cracks and can bottom cracks is increased.
[0023]
After the final cold rolling, finish annealing is performed as necessary.
This finish annealing recovers the processed structure and improves cupping moldability and can bottom moldability. The reason why the finish annealing is performed at a temperature of 100 to 150 ° C. is that the effect is not sufficiently obtained when the temperature is less than 100 ° C. When the temperature exceeds 150 ° C., the solid solution element is excessively precipitated and the DI moldability and the flange moldability are high. It begins to decline. The most desirable finish annealing condition is 115-150 ° C. The finish annealing time is 8 hours or less, particularly preferably 1 to 4 hours.
[0024]
【Example】
Hereinafter, the present invention will be described in detail with reference to examples.
(Example 1)
Al alloy compositions A to E of the present invention composition shown in Table 1 and Al alloys F to M having a composition different from that of the present invention as a comparative example are melt-cast by a conventional method to obtain an ingot (slab) having a thickness of 500 mm. It was. Next, this ingot is chamfered to a thickness of 490 mm, then homogenized at 600 ° C. for 6 hours, then allowed to cool to 350 ° C., then cooled to room temperature, and then the rate of temperature increase is 50 ° C./hour. Then, it was reheated to 520 ° C. and hot rough rolling was performed. The hot rough rolling was performed under the conditions of an end plate thickness of 25 mm, a final pass reduction of 25%, a rolling speed of 120 m / min, and a rolling end temperature of 360 ° C.
Hot finishing rolling was started 180 seconds after the hot rough rolling was completed, and a hot-rolled sheet having a thickness of 2.2 mm was obtained. Hot finish rolling uses a four-stand finish rolling mill, with a total rolling reduction of 90.4%, rolling reduction at each stand F ₁ : 53%, F ₂ : 47%, F ₃ : 44%, F ₄ : 37 %, The rolling end temperature was 330 ° C. The recrystallization rate was 5% before hot finish rolling and 100% after completion of rolling. The recrystallization rate is the area ratio of recrystallized grains in the hot rolled sheet cross section.
After completion of the hot finish rolling, annealing was performed at 400 ° C. for 0 minute (air cooling immediately after reaching the material 400 ° C.) in a continuous annealing furnace. The heating rate at this time was 850 ° C./min, and the cooling rate was 1000 ° C./min. Subsequently, final cold rolling (final cold reduction ratio 87.5%) was performed to a plate thickness of 0.3 mm by a conventional method, and then finish annealing was performed at 115 ° C. for 2 hours to produce an Al alloy plate for a can body.
[0025]
Each Al alloy plate thus obtained was examined for ear ratio, tensile strength (tensile strength (TS) and 0.2% yield strength (YS)), DI formability, and flange formability.
The ear rate was measured by deeply drawing a 57 mmφ disc with a clearance of 30% using a punch having a diameter of 33 mm and a shoulder R of 2.5 mm.
The tensile strength was also measured after heating at 200 ° C. for 20 minutes (paint baking conditions).
The DI moldability was investigated by molding into a DI can barrel (inner diameter 66 mmφ, sidewall plate thickness 103 μm, sidewall tip plate thickness 165 μm) for carbonated beverages.
Flange formability is determined by trimming and cleaning the molded DI can for 20 minutes at 200 ° C., then performing 4-step necking to reduce the inner diameter d of the opening to 57 mmφ, and finally the apex angle of 90 °. The top of the conical jig is pushed in until cracking occurs, the diameter D of the opening when the crack occurs is measured, and the increase rate P of the diameter of the opening is P = [(D−d) / d ] X100% was calculated and evaluated.
The results are shown in Table 2.
[0026]
[Table 1]

[0027]
[Table 2]

[0028]
As can be seen from Table 2, No. A to E of the present invention products have low ear ratios of 2.5% or less, and the limit increase rate of the diameter in flange molding is as large as 1.5% or more, and flange formability is large. Was good. Moreover, the proof stress (YS) after heating at 200 ° C. for 20 minutes was 250 MPa or more, and the pressure resistance at the bottom of the can was at a level with no problem. The DI moldability was also good.
On the other hand, No. F and No. G of the comparative example products each had a high content of Mg or Mn, so that the tensile strength was increased by heating at 200 ° C. for 20 minutes. Softening due to paint baking was insufficient, and ironing cracks occurred in DI molding. No. H was inferior in flange formability due to a large amount of Si and Cu. No. I had a lower Mg content because of its lower Mg content, and No. J had a lower Mn content, which caused seizure in DI molding. No. K had low Cu and Si contents, so the strength decreased. Nos. L and M have high strength due to a large amount of impurities such as Zn or Cr, and iron cracking occurred in DI molding.
[0029]
(Example 2)
A No. A Al alloy shown in Table 1 was melt-cast by a conventional method to obtain an ingot (slab) having a thickness of 500 mm. Next, this ingot was chamfered to a thickness of 490 mm, and then subjected to homogenization, cooling, heat treatment, hot rough rolling, and hot finish rolling in order to obtain a hot rolled coil. After this hot-rolled coil was cooled to room temperature, it was annealed by box-type annealing or continuous annealing, or was not subjected to annealing, and was subsequently cold-rolled by a conventional method to produce an Al alloy plate for a can body. The conditions of homogenization treatment, hot rolling, annealing, and final cold rolling were variously changed as shown in Tables 3 and 5. Tables 4 and 6 show comparative examples implemented under conditions different from those of the present invention.
[0030]
[Table 3]

[0031]
[Table 4]

[0032]
[Table 5]

[0033]
[Table 6]

[0034]
With respect to each Al alloy plate thus obtained, the ear rate, tensile strength, DI formability, and flange formability were investigated by the same method as in Example 1. The ear rate (%) was measured by deeply drawing a 57 mmφ disc with a clearance of 30% using a punch having a diameter of 33 mm and a shoulder R of 2.5 mm. Tables 3 to 6 show production conditions, and Tables 7 and 8 show measurement results. The evaluation criteria were good: an ear rate of 2.5% or less, a yield strength of 260 MPa or more after heat treatment (200 ° C. × 20 minutes), and a limit increase rate of 15% or more of the diameter in flange molding.
[0035]
[Table 7]

[0036]
[Table 8]

[0037]
As is clear from Table 7, the inventive examples (Nos. 1 to 11) had low ear ratios of 2.5% or less and good flange formability. Further, the strength (proof strength) after the heat treatment corresponding to coating baking was 250 MPa or more, a strength level with no problem in the pressure resistance of the bottom of the can, and DI moldability was also good.
[0038]
On the other hand, as is apparent from Table 8, all of the comparative examples had poor characteristics.
That is, No. 1 of the comparative example had a low homogenization temperature, so that the homogenization was insufficient and the ear rate was high.
In No. 2, the homogenization temperature was high and the ingot surface was swollen, and the surface properties after finish rolling were deteriorated.
In Nos. 3 and 4, the cooling rate or the heating rate was slow, so the number density of the precipitates increased and the ear ratio exceeded the reference value.
In No. 5, since the temperature rise was low, the number density of precipitates increased and the ear rate exceeded the reference value.
In Nos. 6 and 7, the final pass R of hot rough rolling, the time t from the end of hot rough rolling to the start of hot finish rolling deviates from the specified value of the present invention, and from the end of rough rolling to the start of finish rolling. Since the recrystallization rate exceeded 30%, the accumulation of distortion was insufficient and the ear rate exceeded the reference value.
In No. 8, since the final plate thickness was thick, the final cold reduction ratio was high. As a result, drawing cracks occurred in DI molding, and the ear ratio exceeded the reference value.
No. 9 had a small finished plate thickness, and seizure occurred after hot finish rolling, and scratches occurred on the surface of the can when it was formed into a can.
In No. 10, since the finish temperature of hot finish rolling was too low, the recrystallization rate after completion was low, and the ear rate exceeded the standard value.
Nos. 11 and 12 have a pass reduction ratio of less than 30% in hot finish rolling and a total reduction ratio of less than 80% in hot finish rolling. Exceeded the reference value.
In No.13, the final cold rolling reduction was high, so that the drawing crack occurred in DI molding, and the ear rate exceeded the standard value.
In No.14, because the finish annealing temperature was high, the strength (proof strength before baking) increased due to precipitation, and ironing cracks occurred.
In No.15, the intermediate annealing temperature was higher than the conditions of the present invention, so the tensile strength by baking at 200 ° C for 20 minutes was improved. The marginal increase rate of the bore diameter was small and the flange formability was inferior.
[0039]
【The invention's effect】
As described above, according to the present invention, an Al alloy sheet for a can body having a low ear ratio and excellent strength, ironing workability, and flange formability after baking is obtained, which has a significant industrial effect. .

Claims

Mg: 0.8-1.4 wt%, Mn: 0.7-1.3 wt%, Fe: 0.2-0.5 wt%, Si: 0.1-0.5 wt%, Cu: 0.1 An aluminum alloy ingot containing 0.3 wt%, Ti 0.005 to 0.05 wt% alone or together with B 0.0001 to 0.01 wt%, with the balance being Al and unavoidable impurities at a temperature of 560 to 620 ° C. Apply homogenization treatment for 1 hour or more in the range, and then perform hot rough rolling by cooling from the homogenization treatment temperature to 450 to 550 ° C. at a cooling rate of 20 ° C./hour or more. After cooling to 30 ° C./hour or higher, it is reheated to 450 to 550 ° C. and subjected to hot rough rolling, and the hot rough rolling is finished with a plate thickness of 12 to 50 mm and an end temperature of 300 to 450. ° C, final pass reduction ratio R ≦ 70−0.2S (S: rolling Subjected under the conditions of the degrees m / min), rough hot rolling finished after t seconds ^{(t = 2.8 × 10 4 exp} (-0.012T), T: hot rough rolling finishing temperature ° C.) within hot finish to Rolling is started, and the hot finish rolling is performed using a tandem hot finish rolling mill having three or more stands, the total reduction rate is 80% or more, the reduction rate at each stand is 30% or more, and the finished sheet thickness is 1.6 to It is applied under conditions of 3.0 mm and end temperature of 290 ° C. or higher, and after hot finish rolling, it is cooled to room temperature, and then kept at 300 to 450 ° C. for 30 minutes or more using a box annealing furnace, or is annealed continuously. Using a furnace, annealing is performed at a temperature rising rate of 100 ° C./min or more at a temperature of 360 to 560 ° C., and immediately after reaching the temperature or after holding for 120 seconds or less, at a cooling rate of 100 ° C./min or more. After cooling to 70 ° C. or lower and annealing, a final cold rolling with a rolling reduction of 60 to 90% is applied, or Aluminum for can body with low ear rate, characterized by subjecting to final cold rolling with a reduction rate of 60 to 90% without annealing after hot finish rolling, and then subjecting to finish annealing at a temperature of 100 to 150 ° C. if necessary Manufacturing method of alloy plate.