JP4094244B2

JP4094244B2 - Titanium for copper foil production drum excellent in surface layer structure and production method thereof

Info

Publication number: JP4094244B2
Application number: JP2001085926A
Authority: JP
Inventors: 卓嗣進藤; 広明大塚; 由尚河原; 満男石井; 義人山下
Original assignee: Nippon Steel Corp
Current assignee: Nippon Steel Corp
Priority date: 2001-03-23
Filing date: 2001-03-23
Publication date: 2008-06-04
Anticipated expiration: 2021-03-23
Also published as: JP2002285267A

Description

【０００１】
【発明の属する技術分野】
本発明は、表層部組織に優れた銅箔製造ドラム用チタンおよびその製造方法に関する。
【０００２】
【従来の技術】
電子部品に組み込まれて使用されているプリント配線板は、銅箔と絶縁基板を張り合わせ、表面に導体パターンをプリントし、不要部分をエッチングで除去して作られている。このプリント配線板には主に電解銅箔が用いられるが、そのような電解銅箔は、高品位の銅原料を硫酸溶液に溶解し、鉛などの不溶性金属を陽極とし、それに対峙させた回転ドラムを陰極として電気化学的にドラム上に銅箔を電着させ、これを連続的に回転ドラムより剥離し、ロール状に巻き取って生産されている。陰極の材料は、耐食性と電着金属の可剥離性から近年チタンが多用されるようになってきた。
【０００３】
銅箔がプリント配線に使用される場合、銅箔の表面粗度はエッチングによって形成される配線パタ−ン（線幅：０．１〜０．５mm）の精度を左右する重要な因子である。この銅箔表面粗度は、銅箔が電着するチタン製陰極ドラムの研磨・整面された粗さを直接に継承する。
【０００４】
また、研磨・整面された陰極ドラムが電解液中において徐々に腐食を受けて、あたかも金属組織検査で用いるエッチングを受けた状態になることがある。このような場合、ドラム表面における金属組織マクロ模様が銅箔上に転写されて仕上げ面粗さの均一性が損なわれ、プリント配線のエッチング不良を生ずるという問題がある。この金属組織の転写模様は、その形態からスクラッチ疵、シラクモ、シジミ等と呼ばれている（以下、スクラッチ疵と総称する。）。
さらにドラム表面は、使用中に電気スパークなどにより表面が荒れてくるため何度も研磨・整面が行われる。その結果、ドラム表面は少しずつ研削されて新しい表面がドラム表面となる。従ってドラム素材には、板面における均一性の他に板厚方向における均一性が要求される。
【０００５】
スクラッチ疵の発生原因は、ドラム素材におけるいわゆるマクロ的不均一組織に起因することがその疵の形態から経験的に知られており、ドラム製造用素材の原素材、すなわちスラブや熱延粗片などに存在するマクロ組織を均一化する試みが行われてきた。例えば、インゴットの鍛造条件あるいは分塊熱延条件、分塊スラブの高温熱処理、熱延厚板の高温長時間焼鈍、などの「加工熱処理法」の検討が行われて、それなりの改善効果が得られてきた。しかしながら、銅箔への品質要求は高度化する一方であり、金属組織の緻密度において、もはや既存の方法ではその要求を満足させることが出来ない状況にある。
【０００６】
従来、一般にチタン展伸材と呼ばれる板材においては、ＪＩＳ規格などに規定される引張り強度、伸びなどの機械的性質と結晶粒度が材料性能特性として求められてきた。しかし、銅箔製造用ドラム素材のように、原素材（インゴット）から加工工程で引き継いできたマクロ的不均一組織の残存を回避するなどのことが求められたり、表面層部位におけるマクロ組織の均一緻密度を求められる場合は無かった。言い換えれば、表面層部位におけるマクロ組織が不均一で緻密度に欠けていても、前記の規格を満足していれば銅箔製造用ドラム以外の用途では特に問題は起こらなかった。
【０００７】
一方、銅箔の製造工程ではチタン製ドラムが硫酸銅水溶液中で回転しながら、その表面に銅を電着させる。その際、ドラム表面は硫酸銅液中で腐食作用を受けて、格子欠陥密度の高い結晶粒界や第二相界面、およびコロニーと呼ばれる近接方位集団粒の境界部が優先的に侵食される。このようにして侵食を受けたマクロ的なスケールにおける不均一模様が銅箔の表面に転写される。この場合のコロニー内では、腐食作用で生ずるエッチピットファセットの形状や大きさ、方向が揃っていて、それらが銅箔に転写されると、極めて軽微な程度のものから明瞭な程度のまでの種々のまだら模様を呈し、これが前述のスクラッチ疵と呼ばれる銅箔の品質不良を引き起こす原因になっていると考えられる。
【０００８】
従来例として、特開昭６０−９８６６号公報においては、スクラッチ疵の解消をはかるために、分塊鍛造および粗熱延における加熱温度を９５０℃以上にすると共に、仕上げ熱延時において圧延板を再度７００℃以下に加熱し、その後、粗熱延時の圧延方向と直交する方向に仕上げ熱延を行うクロス熱延を実施する「β域−α域二段加熱クロス圧延法」によって、均一微細なマクロ模様を有するチタン及びチタン合金板を製造する方法が開示されている。
これによれば、不均一なマクロ模様（以下マクロ不良）の原因は、（ａ）鋳造組織残留によるマクロ不良、（ｂ）変態組織残留によるマクロ不良、（ｃ）粗大結晶粒残留によるマクロ不良、（ｄ）圧延によって生ずる圧延方向に平行なバンド状のマクロ不良、の４タイプに分類され、上記の「β域−α域二段加熱クロス圧延法」は、（ａ），（ｂ），（ｃ），（ｄ）の全てを解消するために有効な方法であるとしている。
【０００９】
しかしながら同法は、粗圧延後の圧延途中に圧延板を再度、仕上げ熱延用の加熱炉に装入する必要があり、これは、加熱温度の異なる二基以上の加熱炉を設置することを前提とするので、生産設備ミル構成上の大きな制約を与える。加えて同法は、仕上熱延を７００℃以下の低温域で圧延で行うために圧延変形抵抗が高くなることから板形状制御が難しく、また、いきおい再結晶温度以下の低温域において圧延を行うことが多くなり、前述のタイプ（ｄ）のマクロ不良が発生する傾向が増大するなどの難点を有している。
【００１０】
一方、特開平９−１７６８０９号公報においては、▲１▼「α域一段加熱クロス圧延法」すなわち分塊鍛造あるいは分塊圧延にて得たスラブを、粗熱延、仕上げ熱延を行うに際して、粗熱延時のスラブ加熱温度を７００℃以上からβ変態温度未満の温度範囲にすると共に、直交する長手方向と幅方向に対して圧延を行い、且つ長手方向への圧延による加工真歪み量をε_L、幅方向への圧延による加工真歪み量をε_wとするとき、両加工真歪み量の間にε_L×ε_w≧０．２５が成り立つ条件で圧下することを特徴とする、不均一マクロ模様のないチタンまたはチタン合金板の製造方法が開示されている。
さらに、▲２▼粗熱延に供するスラブとして、β変態温度未満で熱間加工が施された後、β変態温度以上に加熱することにより結晶粒が微細化されたスラブを用いること、あるいは▲３▼熱延板を８００℃以下で焼鈍した後、冷延および焼鈍を付加して平均結晶粒径が５０μｍ以下とする製造方法が開示されている。
【００１１】
しかしながら、これらのうち、▲１▼はクロス圧延における板長手方向と板幅方向の総圧下量を規定しているが、クロス圧延を実施すべき適切な圧延温度範囲を規定しておらず、技術的に極めて不十分である。また▲２▼は、β単相域にスラブを再加熱すると粒成長が顕著に起こり、組織の微細化は達成されない。さらに▲３▼は、熱延板を焼鈍する温度と、次いで行う冷間圧延の圧延率と最終焼鈍の温度、保定時間の詳細が明示されておらず、平均粒径が５０μｍのチタン板が得られるべき技術要件の規定が不十分である。
【００１２】
【発明が解決しようとする課題】
本発明の課題は、前記の事情を考慮して、板面表層部において腐食後にスクラッチ疵の原因となるマクロ的不均一組織を示さない、表層部組織に優れた銅箔製造ドラム用チタンおよびその製造方法を提供するものである。
特に製造方法に関しては、熱延において前記のような設備上の制約あるいは製造上の難点を有するいわゆる二段加熱圧延法によらない、いわゆる一段加熱圧延法と、それを含む一貫製造方法を提供するものである。
【００１３】
【課題を解決するための手段】
発明者らは、銅箔製造用ドラムに供するチタンの表層部組織とスクラッチ疵等との関係を鋭意検討の結果、表層部組織の満たすべき条件を見出し、このための成分や、設備上の制約の少ない製造方法を限定するに至って本発明を完成させたもので、その要旨とするところは以下の通りである。
【００１４】
（１） mass％で、
Ｆｅ：０．０４％以下、Ｏ：０．１％以下
を含有し、残部チタンおよび不可避不純物からなり、最表面から１／３板厚にわたる板面に平行な表層部において、α相の平均粒径が４０μｍ以下であり、さらに、α相集合組織が［０００１］０°〜±４５°ＴＤ、かつ、［０００１］０°〜±２５°ＲＤ（ただしＴＤは板幅方向、ＲＤは圧延方向）であることを特徴とする表層部組織に優れた銅箔製造ドラム用チタン。
【００１５】
（２）前記（１）に記載の銅箔製造ドラム用チタンを製造するに際し、電子ビーム溶解法によって溶解鋳造した厚さ３００mm以上の矩形断面スラブを、β域に加熱し、β域において圧下比３以上の分塊圧延あるいは分塊鍛造を行いβ相再結晶組織を形成させた後、直ちに、β域加工終了温度〜７００℃の範囲を冷却速度２００℃／ｈｒ以上で冷却し、その後８８０℃以下に加熱して粗熱延を行い、該粗熱延後再加熱することなく、粗熱延の圧延方向と直交する方向に圧延するクロス熱延をクロス圧延比が１／１０〜６／１０となるようにした仕上げ熱延を、６５０〜７５０℃の温度範囲で行い、さらに５７０〜６８０℃の温度範囲において１〜１０ｈｒの焼鈍を行うことを特徴とする表層部組織に優れた銅箔製造ドラム用チタンの製造方法。
（３）前記焼鈍後、さらに圧延率が２５〜５０％の冷延を行い、その後６００〜７１０℃の温度範囲で１０〜３０ｍｉｎの焼鈍を行うことを特徴とする前記（２）に記載の表層部組織に優れた銅箔製造ドラム用チタンの製造方法。
【００１６】
【発明の実施の形態】
以下、本発明を詳細に説明する。
発明者らは、スクラッチ疵の生成機構、特にその冶金学的原因を解明するべく鋭意研究を行った結果、まず、スクラッチ疵の直接原因となる熱間圧延板中のマクロ的不均一組織の組織学および結晶学との関連性について以下の重要な知見を得た。すなわち発明者らは、純チタンの厚板圧延等における熱間圧延集合組織形成に関する詳細な研究を行った。
【００１７】
その結果、
▲１▼ スラブの加熱温度をβ域とし、β〜α域において連続的にスラブ長手方向と平行方向の一方向（もしくはリバース方向）圧延を施すβ域加熱β〜α域連続圧延材は、［０００１］軸が板幅方向に９０°傾く、いわゆるTransverse-textureを示す。これに対して、
▲２▼ スラブの加熱温度をα域とし、α域において連続的に一方向（もしくはリバ−ス方向）圧延を行うと、［０００１］軸が板幅方向に３５〜４５°傾く、いわゆるSplit-TD-textureを示す。さらに、
▲３▼ α域の低温域（約６００〜約７５０℃）において、圧延方向を９０°回転させるクロス圧延を行うと、［０００１］軸が板法線方向とほぼ平行（±２５°程度以内）になる、いわゆるCenter-Pole-texture を形成する。
ことを明らかにした。
このような種々の熱間圧延を行ったチタン板において、前記▲１▼、▲２▼の場合にはマクロ的不均一組織が生成する傾向が高いが、▲３▼の方法で製造したチタン板は、細粒でマクロ的不均一組織が極めて軽微な程度にまで減少することを見出した。
【００１８】
前記▲３▼の方法によって製造したチタン板において、マクロ的不均一組織が減少する冶金学的理由は、以下のように考えられる。
すなわち、一般に一方向（もしくはリバ−ス）圧延を行うと、圧延面に対して約４５°の傾角を示す剪断すべり変形帯が形成されるが、クロス圧延を行うと圧延面法線軸に関してさらに９０°回転した４５°剪断すべり変形帯が付加的に導入される。α域の低温域における圧延では圧延双晶が形成されるので、単一のα相結晶粒に注目すると、結晶粒内はクロス圧延によって生ずる交叉的なすべりと双晶変形とによって、組織が極めて微細に分断される。このようなすべり変形に基づく加工歪みは再結晶駆動力を増大させ圧延中の再結晶を促進し、均一細粒組織を形成する。また高転位密度を有する双晶も擬似的な粒界を構成するので、これも組織の細粒化を促進する。
【００１９】
このようにα域の低温域においてクロス圧延を行えば、マクロ的不均一組織の破壊と細粒化が促進され、スクラッチ疵の低減にとっては極めて有効な方法となる。言い換えれば、チタン板のマクロ的不均一組織を解消するためには、熱間仕上げ圧延段階においてCenter-Pole-texture を形成するように圧延を行うことが極めて重要な要件であることを示している。
【００２０】
前述のようにマクロ的不均一組織を解消するためには、粗熱延から仕上げ熱延工程においてクロス圧延を行い、（ｃ）粗大結晶粒残留を生じさせないことが極めて重要となるが、その効果を一層有効に発現させるためには、粗熱延以前のスラブ組織が、（ａ）鋳造組織残留によるマクロ不良や、（ｂ）変態組織残留によるマクロ不良を生ずるような粗大な不均一組織を示してはならない。
【００２１】
発明者らは、粗熱延に供するスラブの形状、スラブ分塊圧条件、スラブ組織等と熱間圧延板中のマクロ的不均一組織形成との関係について詳細に研究し、電子ビ−ム溶解法によって溶解鋳造したスラブ厚さが３００mm以上である厚肉矩形断面チタンスラブを、β域に加熱しβ域において圧下比３以上の分塊圧延あるいは分塊鍛造を行いβ相再結晶組織を形成させた後、直ちに、β域加工終了温度〜７００℃の範囲を冷却速度２００℃／ｈｒ以上で冷却を行いβ→α変態に基づく微細な変態組織を形成させると、熱間圧延板のマクロ的不均一組織が減少することを知見した。
【００２２】
消耗電極式アーク溶解（ＶＡＲ）法によって製造した鋳塊は、通常、断面直径が約５００〜８００mmφ程度の円柱形状を示すので、この円柱状鋳塊を熱間圧延するためには、分塊圧延もしくは分塊鍛造によって矩形断面スラブに加工する必要がある。その場合、円柱状鋳塊の円柱軸方向と直交する方向（外周面の法線方向）から圧延もしくは鍛造による圧縮加工を施すが、このとき加工後のスラブ断面幅方向の端部に相当する加工前円柱状鋳塊断面の外周部付近は、鋳塊中心部付近に比べて充分な圧下比を確保することが困難になる。その結果、分塊圧延・鍛造後のスラブ断面組織が全体として不均一な組織を示すと言う難点を有する。
【００２３】
ＶＡＲ鋳塊の分塊圧延・鍛造は、一般的にβ域加熱されβ域からα域にわたって加工される。これはα域加熱α域加工を行うと、素材がβ→α変態に基づく組織微細化作用を被る機会がなくなり、結果として（ａ）鋳造組織残留によるマクロ不良の発生を回避できないためである。このように、ＶＡＲ円柱状鋳塊の外周部は低い圧下比の加工を余技なくされ、同部位には粗大なβ相未再結晶組織を呈しやすく、（ｂ）変態組織残留によるマクロ不良も発生する傾向が大きい。
【００２４】
これに対して本発明においては、電子ビーム熔解（ＥＢＲ）法によりハース溶解後に矩形スラブ鋳型に鋳造する矩形断面スラブを製造するので、前述のようなＶＡＲ円柱状鋳塊特有の分塊圧延・鍛造後のスラブ断面組織が不均一になるという難点は生じない。
【００２５】
また、鋳塊矩形断面における板厚を３００mm以上、望ましくは６００mm程度の厚肉にして製造するので、仮に分塊圧延スラブの圧延仕上げ板厚を１００mmとすると、β域加熱後のβ域圧延における圧下比が３（加工率６６％）以上から最大６（加工率８３％）程度まで確保することが可能である。また鋳塊の板厚が充分に厚いので、圧延中の温度がα域にまで低下することなく、完全なβ域における圧延が実施される。
このように、β域において圧下比が３以上の加工を行うと、加工中にβ相が再結晶を生じて微細化するうえ、その後の冷却過程におけるβ→α変態によってさらに微細な変態β組織（微細αラメラー組織）が得られる。これによって、熱延工程に供するスラブ組織としては充分な組織微細化と結晶方位のランダム化が実現される。
【００２６】
また、スラブ冷却過程において、α域の粒成長が顕著に生ずるようなβ域加工終了温度から７００℃からの温度範囲に長時間滞在すると、αラメラー組織の粗大化が起こり望ましくないので、このような温度範囲はできるだけ急速に冷却される必要がある。
発明者らは、特にβ域加工終了温度〜７００℃の範囲を冷却速度２００℃／ｈｒ以上で冷却を行い、β→α変態に基づく微細な変態組織を形成させることができることを見出した。これに対して、冷却速度が２００℃／ｈｒ未満の緩冷却を行う場合には、変態組織の粗粒化が起こってマクロ不良の原因となり望ましくない。
【００２７】
さらに本発明においては、前述の方法によって鋳造、分塊圧延もしくは分塊鍛造し、冷却したチタンスラブの熱延において、粗熱延における加熱温度を８８０℃以下とし、粗熱延後再加熱することなく連続して６５０〜７５０℃で仕上げ熱延を行う製造方法を採ることが好ましい。
【００２８】
この仕上げ熱延においては、粗熱延の圧延方向と直交する方向に圧延するクロス熱延を取り入れ、クロス圧延比が１／１０〜６／１０となるように圧延する。なお、クロス圧延比＝クロス圧延前の前段総圧下比／クロス後の後段総圧下比である。
この場合、粗熱延におけるスラブ加熱温度を８８０℃以下とする理由は、β変態温度に相当する８８０℃以上に加熱すると、微細αラメラー組織がβ相に変態しβ相の粒成長を起こすので、これを避ける必要があるためである。
【００２９】
発明者らは、６００mm板厚のＥＢＲ溶解鋳塊をβ域加熱分塊圧延、β域仕上げ後急冷して製造した、寸法が１５０mm（ｔ）×１２６０mm（ｗ）×８５０mm（Ｌ）のスラブ（β分塊圧延圧下比：４．０）から、寸法が１５０mm（ｔ）×１７０mm（ｗ）×２５０mm（Ｌ）の直方体圧延用ブロックを１０数個採取し、これらの圧延用ブロックをα域（８５０℃）に加熱し、以下の３種類の圧延実験を行い、詳細な圧延条件を検討した。
【００３０】
（Ａ）：α域一段加熱連続一方向圧延（後段圧延温度制御）
α域一段加熱連続一方向圧延においては、前段粗圧延を１５０mm厚から３０mm厚まで７パスで圧延し、その後、後段仕上げ圧延開始温度を５００〜７５０℃の範囲で変化させ、連続してブロック長手方向（Ｌ方向）に３０mm厚から９．０mm厚まで４パスの圧延を施した。このときの前段圧延は圧延率８０％（圧下比５．０）、後段圧延は圧延率７０％（圧下比３．３）である。また圧延開始から圧延終了までの全圧延率は９４％（圧下比１６．７）である。図１（観察倍率×１００）は、圧延材の表面下１mm板厚層部位板面平行面のミクロ組織観察結果である。
【００３１】
図１の（１）、（２）に示した、後段圧延の圧延開始温度が７５０℃および７００℃である高温圧延材の場合、微細な再結晶粒が生成しているが、後段圧延温度が（３）６００℃、（４）５００℃と低下するに従い、圧延双晶を含む未再結晶混粒組織を示している。マクロエッチ組織観察においても、特に低温後段圧延圧延材（６００℃、５００℃）に、圧延方向（Ｌ方向）に平行方向に延伸したマクロ粒が顕著に発生していることが確認された。従って、ミクロ組織観察とマクロ組織観察結果から、α域一段加熱連続一方向圧延においては圧延方向平行な顕著なマクロ的不均一組織を形成しやすく、その傾向は後段圧延開始温度の低下と共に顕著になることが明らかである。
【００３２】
（Ｂ）：α域一段加熱連続クロス圧延（後段クロス圧延温度制御）
α域一段加熱連続クロス圧延においては、前段粗圧延を１７０mm（Ｗ）方向に１５０mm厚から８４mm厚まで４パスでＣ方向圧延し、その後冷却ならびに再加熱することなく、連続して直ちに９０度回転し、後段圧延開始温度を５００〜７５０℃の範囲で変化させ、Ｌ方向に８４mm厚から９．０mm厚まで８パスのクロス圧延を施した。このときの前段粗圧延は圧延率４４％（圧下比１．７８）、後段仕上げ圧延は圧延率８９．２％（圧下比９．３３）である。従って、クロス比Ｃ／Ｌ＝１／５．２４、また圧延開始から圧延終了までの全圧延率は９４％（圧下比１６．７）である。
【００３３】
図２（観察倍率×１００）は、圧延材の表面下１mm板厚層部位板面平行面のミクロ組織観察結果である。図２に見られるように、（１）ＥＢ５，後段圧延開始温度７５０℃、（２）ＥＢ６，後段圧延開始温度７００℃、材においてはいずれも微細な再結晶等軸組織を示し、極めて良好な組織となる。（３）ＥＢ７，後段圧延開始温度６００℃材は、僅かに圧延方向に平行な未再結晶延伸粒が観察される。
一方、マクロエッチ組織観察においても、ＥＢ５、ＥＢ６材は未再結晶α延伸粒を含む不均一マクロ組織を示すことなく、極めて均一なマクロ組織を示していた。これに対してＥＢ７材のマクロ組織は、僅かながら未再結晶α延伸粒を含むマクロ的不均一組織を呈していた。このようにα域一段加熱連続クロス圧延においては、後段圧延開始温度を７００〜７５０℃とすると、極めて均一なマクロ組織が得られ、α域一段加熱連続一方向圧延に比べて顕著な改善効果が見られる。
【００３４】
（Ｃ）：α域二段加熱クロス圧延（後段クロス圧延温度制御）
α域二段加熱クロス圧延においては、前段粗圧延を１５０mm厚から６０mm厚まで（ＥＢ９）あるいは３５mm厚まで（ＥＢ１０）、５〜７パスでＣ方向圧延し水冷した。その後圧延板を切断し、７００℃に再加熱し、前段粗圧延の圧延方向から９０度圧延方向を回転させるＬ延べクロス圧延を行った。その際、後段仕上げ圧延の圧延開始温度は７００℃一定として、６０mm厚あるいは３５mm厚から９．０mm厚までＬ方向に５〜７パスのクロス圧延を施した。
このときの前段圧延は圧延率６０％（圧下比２．５）あるいは圧延率７７％（圧下比４．２８）、後段圧延は圧延率８９．２％（圧下比９．３３）あるいは圧延率７４．２％（圧下比３．８）である。従って、クロス比はＥＢ９はＣ／Ｌ＝１／２．６６、あるいはＥＢ１０はＣ／Ｌ＝１．０／０．９０である。また圧延開始から圧延終了までの全圧延率は９４％（圧下比１６．７）である。
【００３５】
図３（観察倍率ｘ１５）、および図４（観察倍率×１００）は、圧延材の表面下１mm板厚層部位板面平行面のミクロ組織観察結果である。図において、（１）ＥＢ９，Ｃ／Ｌ＝１．０／２．６６においてはマクロ的延伸粒がほぼ解消されており、細粒均一組織が形成されている。図２に示した再加熱を行うことなく７００℃において連続クロス圧延を行ったＥＢ６（クロス比Ｃ／Ｌ＝１．０／５．２４）と、図４に示した７００℃再加熱クロス圧延を行ったＥＢ９（クロス比Ｃ／Ｌ＝１．０／２．６６）を比較すると、両者はほぼ良好な微細な均一等軸組織を示しており、この温度域においてクロス圧延を行うことは、マクロ的不均一組織の解消に極めて有効であることが確認された。
【００３６】
これに対してクロス比を強化し、Ｃ／Ｌ＝１．０／０．９０とした図３（２）および図４（２）に示したＥＢ１０材は、α延伸粒を形成して、マクロ的不均一組織を呈している。この理由としては、ＥＢ１０材はクロス比は１．０／１．０に近い条件となったが、後段の圧下比が約３．８と低い圧延となったために、α延伸粒の解消が不十分であったものと考えられる。一方、さらに観察面積を拡大したマクロエッチ組織観察においては、ＥＢ９、ＥＢ１０材は未再結晶α延伸粒を含む不均一マクロ組織をところどころに示していた。
【００３７】
以上述べた（Ａ），（Ｂ），（Ｃ）の結果から、電子ビーム熔解によって鋳造した矩形断面スラブを、前述のごとくβ域において分塊圧延（もしくは鍛造）を施し冷却したチタンスラブに、粗熱延及び仕上げ熱延を順次連続的に施す工程において、粗熱延における加熱温度を８８０℃以下とし、粗熱延後再加熱することなく連続して行う仕上げ熱延において、粗熱延との圧延方向が直交するクロス熱延を、温度範囲が６５０℃〜７５０℃で、クロス圧延比（＝クロス圧延前の前段総圧下比／クロス後の後段総圧下比）が２／１０（約１．０／５．２４）〜４／１０（約１．０／２．６６）となるように圧延することが最も適切であることが判明した。この場合、クロス圧延比の上限は６／１０までに、下限は１／１０にまで拡張しても、マクロ的不均一組織は解消可能であった。
【００３８】
前記の方法で製造したチタン厚板（熱間圧延まま材）は、平坦度などの形状が不良である場合が多いので、５７０〜６８０℃の温度範囲で焼鈍を行う。焼鈍は、真空クリープ矯正焼鈍設備（ＶＣＦ：Vacuum Creep Flattening ）を用いて、形状矯正を兼ねた焼鈍処理を行うことが好ましい。
表１に厚板圧延まま材、各種ＶＣＦ焼鈍温度とマクロ組織の関係を、および図５にＶＣＦ焼鈍温度とＶＣＦ焼鈍材のミクロ組織における平均結晶粒径の関係を示す。
【００３９】
【表１】

【００４０】
表１から明らかなように、圧延まま材にはマクロ的不均一が残留したが、５７０℃以上の温度において焼鈍を行うと、再結晶が完了してマクロ的不均一組織が完全に消失する。ＶＣＦ焼鈍の保定温度における保定時間は１〜１０時間の範囲で行うが、図５に示すようにＶＣＦ焼鈍材の平均結晶粒径が４０μｍ以上の粗粒になると、ＶＣＦ焼鈍材をドラム素材として供する場合はドラム研磨性が不良になる。
【００４１】
また、ＶＣＦ焼鈍に続いて冷間圧延工程および焼鈍工程を付加して銅箔ドラム用チタン厚板を製造する場合においては、冷間圧延前素材の平均結晶粒径は、過剰に粗大であるとマクロ的不均一組織の原因にもなるなどの難点を生じるので、ＶＣＦ焼鈍まま材の平均粒径は実質的に１０〜７０μｍの範囲となることが適当である。
【００４２】
図５に示すように、ＶＣＦ焼鈍温度範囲を５７０〜６８０℃として、保定時間範囲を１〜１０ｈｒにすることによって、ＶＣＦ焼鈍まま材においても、また後述するようにＶＣＦ焼鈍に続いて冷間圧延工程および焼鈍工程を付加して銅箔ドラム用チタン厚板を製造する場合においても、細粒均一なマクロ的不均一模様のない電解銅箔製造のための電着ドラム用チタン厚板を提供可能にする。なお、ＶＣＦでない通常の焼鈍による場合も、温度や時間は上記の条件によって行う。
【００４３】
α域一段加熱連続クロス圧延を行ったチタン厚板圧延板に、６３０℃において５ｈｒ保定を行うＶＣＦ焼鈍を施し、全圧延率が約４０％となる冷間圧延を加えた後に、６００〜７５０℃の温度範囲において１０〜３０ｍｉｎの保定を行う焼鈍を付加した。図６は、そのような焼鈍板の板厚最表層部位から１／３板厚相部位における断面金属ミクロ組織から観察したα相の平均結晶粒径と焼鈍温度、焼鈍保定時間の関係を表す図である。
【００４４】
図６中で網掛けハッチで囲った領域は、本発明における最終焼鈍の適切な温度、保定条件範囲である。例えば最終焼鈍チタン板の平均結晶粒径が５μｍ以上であれば、マクロ不均一組織はほぼ解消されている。また平均結晶粒径が４０μｍ超であると、マクロ的不均一組織は解消されているが、粗大な結晶粒径が銅箔表面の粗度を過剰に大きくし、ドラム研磨性が不良になるなどの難点を生じて望ましくない。従って、平均結晶粒径が５〜４０μｍの範囲内であれば良いが、これを実現する保定時間としては、例えば１０ｍｉｎ以下であれば板面内の均一な温度分布が達成されず、結晶粒径が混粒化するのでよくない。
【００４５】
また、保定時間が３０ｍｉｎ以上であれば表面酸化が過剰になり、後続の脱酸化膜工程の負荷が大きくなり、また生産性が悪いなどの理由でよくない。
以上の理由から、冷間圧延後の最終焼鈍条件としては、６００〜７１０℃の温度範囲において１０〜３０ｍｉｎの保定を行うように限定した。また最終焼鈍以前の冷間圧延率に関しては、ＶＣＦ焼鈍を行ったチタン厚板に、厚板長辺方向に全圧延率が２５〜５０％の範囲で行うことが望ましい。圧延率は２５％未満であれば、最終焼鈍板の金属組織は混粒化しやすく、マクロ不均一組織が残存する場合があるため望ましくない。圧延率は５０％以上であれば、圧延板形状が不良になるのでこれも望ましくない。
【００４６】
次に、前記のような本発明に基づく一連の製造方法によって製造されるチタン厚板は、適切な温度域において適切なクロス圧延比に基づいてクロス熱間圧延を実施することから、圧延もしくは焼鈍板において得られる集合組織は、例えば熱延まま板においてCenter-Pole-texture ，Split-TD-texture，およびSplit-RD-texture，Transverse-textureが複合的に混在することが特徴である。
【００４７】
これが原因でＶＣＦ焼鈍まま、あるいはＶＣＦ焼鈍＋冷間圧延＋最終焼鈍においては、Split-TD-texture（＋Split-RD-texture）が形成されやすくなる。従って、チタン素材に施される加工熱処理工程がどのような場合であっても、板厚最表層部位から１／３板厚層部位における集合組織は、Center-Pole-Texture ，Split-TD-Texture および Split-RD-textureの単一もしくは複合的集合組織、すなわち［０００１］０°〜±４５°ＴＤ、かつ［０００１］０°〜±２５°ＲＤ（ただしＴＤは板幅方向、ＲＤは圧延方向）が主方位となる集合組織を有する上に、表面〜１／３板厚部位における表面層α相粒の平均結晶粒径が４０μｍ以下である、細粒均一なマクロ的不均一模様のない電解銅箔ドラム用チタン厚板であれば良い。
【００４８】
従って、本発明の実施の形態において詳細に述べた方法が、その製造に最も適していることはこれまで述べた通りであるが、それ以外の方法であっても特に制限するものではない。また本発明が適用される純チタンの化学組成については、望ましくは、Ｏ≦０．１mass％、かつＦｅ≦０．０４mass％とする。
【００４９】
Ｏ≦０．１mass％とする理由は、チタン厚板をドラム状のリングに成形し、両端部を合わせ溶接する際に、溶接中に２００ｐｐｍ程度の酸素が付加的に混入（酸素ピックアップ）が生じて、溶接部が母材部よりも硬度が高くなることがあり、Ｏが０．１mass％より多いと、溶接部の硬度増加が一層顕著になって、リング合わせ溶接部に冷間加工と焼鈍を加えその表面部を研磨精面しても、溶接部線に沿って線状の高さ０．５μｍ程度の凸状表面欠陥が残留して、銅箔に凹状の線状欠陥が転写されるためである。
【００５０】
また、Ｆｅ≦０．０４mass％とする理由は、Ｆｅが０．０４mass％より高いと、結晶粒界等に偏在するＦｅを基点にする孔食状のｐｉｔｔｉｎｇが著しくなり、これが銅箔に転写されるので、銅箔に凸状の表面欠陥が形成されるためである。
【００５１】
【実施例】
（実施例１）
本実施例は、本発明請求項１に関わるものである。
Ｏが０．０２〜０．１３mass％、Ｆｅが０．０１３〜０．０６mass％である２０種類の高純度チタンのＥＢＲチタン鋳塊、もしくはＶＡＲ鋳塊を溶解製造し、これらの鋳塊に以下の表２中の▲１▼〜▲６▼に示す加工熱処理を付与し、チタン厚板を製造した。
【００５２】
例えば製造工程▲１▼では、ＥＢＲ鋳塊（寸法：６００mm厚×１３００mm幅×２０００mm長）をβ域（１０００℃）に加熱し、β域において１５０mm厚まで圧下比４．０の分塊圧延を行った。
このときβ域加工終了温度（９００℃）〜７００℃間の冷却速度は２５０℃／ｈｒとした。その後α域（８５０℃）に加熱し、分塊スラブの板幅方向（Ｃ方向）に圧下比３．０の粗熱延を行い、引き続いて仕上げ熱延として粗熱延板圧延方向に対して圧延方向が直交する（Ｌ方向）クロス熱延を、温度範囲が７００〜７５０℃の間において、クロス圧延比がＣ／Ｌ＝６／１０となるように行った。その後、真空クリープ板形状矯正ＶＣＦ焼鈍を（６５０℃×１ｈｒ）行った。
【００５３】
製造工程▲２▼では、ＥＢＲ鋳塊（寸法：６００mm厚×１３００mm幅×２０００mm長）をβ域（１０００℃）に加熱し、β域において１５０mm厚まで圧下比４．０の分塊圧延を行った。このときβ域加工終了温度（９００℃）〜７００℃間の冷却速度は３００℃／ｈｒとした。
その後α域（８５０℃）に加熱し、分塊スラブの板幅方向（Ｃ方向）に圧下比２．０の粗熱延延を行い、引き続いて仕上げ熱延として粗熱延との圧延方向が直交する（Ｌ方向）クロス熱延を、温度範囲が７００〜７５０℃の間においてクロス圧延比がＣ／Ｌ＝４／１０となるように行った。さらに真空クリープ板形状矯正ＶＣＦ焼鈍を（６５０℃×１ｈｒ）行った。その後圧延率４０％の冷間圧延を行い、最終焼鈍（６６０℃×２０ｍｉｎ）を加えた。以下、製造工程▲３▼〜▲６▼については表２に記した方法である。
【００５４】
【表２】

【００５５】
得られたチタン厚板の板面を、表面下１mm層板面および１／３板厚層板面までそれぞれ研磨して、硝フッ酸にて腐食しマクロ模様の目視検査を行った。表３にそれらの評価結果を示す。表３中で結晶粒径は１／３板厚層部位板面の平均結晶粒径を意味し、集合組織は同部位板面層の（０００１）正極点図測定によって得られた主方位（最高強度極密度ピーク位置）、および副方位（第二番目、および第三番目強度の極密度ピーク位置）を表す。
ここで、Split-TD（±４２°）は［０００１］軸がＴＤ方向に±４２°傾くSplit-TD-textureを、Center-Pole は、Split-RD（０°〜±２５°）、およびSplit-TD（０°〜±２５°）が複合的に混在する［０００１］軸がＮＤ（板法線）方向から２５°以内にあるCenter-Pole-Texture を示す。Transverseは、［０００１］軸がＴＤ方向に７０°以上傾くTransverse-textureを意味する。
【００５６】
表３中に示す製造工程▲１▼および▲２▼で製造したサンプル記号Ａ１〜Ａ９の本発明チタン板は、マクロ模様の発生はほとんど見られず、結晶粒径は４０μｍ以下でα相集合組織が［０００１］０°〜±４５°ＴＤ、かつ［０００１］０°〜±２５°ＲＤとなった実施例である。比較例のサンプル記号Ａ１０〜Ａ２０は軽度な、場合によっては重度なマクロ模様が発生し、結晶粒径およびα相の集合組織ともに上記の要件を満たさない例である。
【００５７】
【表３】

【００５８】
（実施例２）
本実施例は、本発明請求項２に関わるものである。
電子ビーム溶解法によって鋳造した、Ｏ成分が０．０４mass％、Ｆｅ成分が０．０３mass％で、スラブ厚さが２５０〜６５０mmの範囲にある矩形断面チタンスラブを、α域（８５０℃）およびβ域（９５０℃、１０００℃）に加熱し、圧下比が１．５〜５．０の範囲で分塊圧延を行った。その後、直ちに７００℃〜β域加終了温度の温度範囲を冷却速度：１００〜５００℃／ｈｒで冷却した。
これらの種々の分塊圧延スラブを実施例１中の「工程▲１▼」と同様の方法によって、厚板圧延、およびＶＣＦ焼鈍を行い、得られたチタン厚板の板面を１／３板厚層部位まで研磨して、硝ふっ酸にて腐食しマクロ模様の目視検査を行った。
【００５９】
また、ＯとＦｅ成分がほぼ同量のＶＡＲ溶解を行った直径が８００mmの円柱状鋳塊（０．０４mass％Ｏ，０．０３１mass％Ｆｅ）を、β域（１０００℃）に加熱し、圧下比（直径部の圧下比）が２．８〜４．５の範囲で分塊圧延を行い、その後直ちに７００℃〜β域ないしα域における加工終了温度の温度範囲を冷却速度：１００〜３００℃／ｈｒで冷却した。これらの種々の分塊圧延スラブを実施例１中の「工程▲１▼」と同様の方法によって、厚板圧延およびＶＣＦ焼鈍を行い、得られたチタン厚板の板面を１／３板厚層部位まで研磨して、硝ふっ酸により腐食しマクロ模様の目視検査を行った。表４にこれらの厚板の製造条件とマクロ組織評価結果をまとめて示す。
表４から明らかなように、本発明実施例Ｂ１〜６はいずれもマクロ組織は消滅（評価○もしくは◎）しているが、比較例Ｂ７〜１２はいずれもマクロ模様がみられ、特に比較例のサンプル記号Ｂ１１およびＢ１２はマクロ模様が顕著であった。
【００６０】
【表４】

【００６１】
（実施例３）
本実施例は、本発明請求項２に関わるものである。
電子ビーム溶解法によって溶解鋳造したＯ成分が０．０３mass％、Ｆｅ成分が０．０２８mass％で、スラブ厚さが６３０mmである厚肉矩形断面チタンスラブを、β域（９８０℃）に加熱し、β域において圧下比４．２の分塊圧延を行いβ相再結晶組織を形成させた後、直ちに７００℃〜β域加工終了温度の範囲を冷却速度２５０℃／ｈｒで冷却を行い、β→α変態に基づく微細な変態組織を形成させた、板厚１５０mmの分塊圧延スラブを複数枚製造した。
これらの分塊圧延スラブから、寸法が１５０mm厚×１３００mm幅×２０００mm長の厚板圧延用素片を採取して、表５にサンプル記号Ｃ１〜Ｃ１１として示す製造条件に基づくチタン厚板を製造した。
【００６２】
チタン厚板圧延工程におけるクロス以前の粗圧延は、分塊圧延スラブの板幅方向（Ｃ方向）に圧下し、圧延方向を９０°変換して（分塊圧延方向と同じ方向：Ｌ方向）、仕上げ圧延を行った。その際、仕上げ圧延開始温度を種々の温度に設定して実施した。
表５から明らかなように、ＶＣＦ焼鈍板において、本発明実施例Ｃ１〜６はいずれもマクロ組織は消滅（評価○もしくは◎）しているが、比較例Ｃ７〜１１はいずれもマクロ模様がみられ、特に比較例のサンプル記号Ｃ１０およびＣ１１はマクロ模様が顕著であった。
【００６３】
【表５】

【００６４】
（実施例４）
本実施例は、本発明請求項３に関わるものである。
電子ビーム溶解法によって鋳造した、Ｏ成分が０．０４mass％、Ｆｅ成分が０．０３mass％で、スラブ厚さが６００mmの範囲にある矩形断面チタンスラブを、β域（９５０℃）に加熱し、圧下比が４．０の分塊圧延を行い、その後直ちにβ域加熱終了温度〜７００℃の温度範囲を冷却速度３００℃／ｈｒで冷却した。
このようにして製造した板厚１５０mmの分塊圧延スラブを複数枚準備して、その後α域（８５０℃）に加熱し、分塊スラブの板幅方向（Ｃ方向）に圧下比２．０の粗熱延延を行い、引き続いて仕上げ熱延として、粗熱延との圧延方向が直交する（Ｌ方向）クロス熱延を、温度範囲が７００〜７５０℃の間においてクロス圧延比がＣ／Ｌ＝４／１０となるように行った。
【００６５】
さらに、真空クリープ板形状矯正ＶＣＦ焼鈍を、温度範囲が５００〜７００℃、保定時間が１〜２０ｈｒとなる条件で行った。詳しいＶＣＦ焼鈍条件は表６中に示すが、得られた焼鈍板の表面下１mm層板面および表面下１／３板厚層板面におけるマクロ組織および結晶粒径を評価した。ＶＣＦ焼鈍板を電解ドラム板として使用する場合は平均結晶粒径が４０μｍ以下に、ＶＣＦ焼鈍板に冷間圧延と最終焼鈍を付加する場合は平均結晶粒径は７０μｍ未満である必要がある。表６中に平均結晶粒径が７０μｍ以上となる場合は粗粒と記した。
【００６６】
【表６】

【００６７】
表６中で、記号Ｄ１〜Ｄ７，Ｄ９はマクロ模様はかすかな程度（評価○）であったが、平均結晶粒径が７０μｍ以上となるので宜しくない。記号Ｄ２１〜Ｄ２８はマクロ模様が残存している（評価△、×）ので宜しくない（圧延組織が残存して平均結晶粒径が測定できない場合は不定と記した）。記号Ｄ８，Ｄ１０〜Ｄ１２，Ｄ１４〜Ｄ２０はマクロ模様もかすかな程度（評価○）で、結晶粒径も７０μｍ未満となる本発明例である。表６中の＊印をつけた発明例は、ＶＣＦ焼鈍＋冷間圧延＋最終焼鈍工程を経てドラム板を製造する場合の例である。
【００６８】
（実施例５）
本実施例は、本発明請求項５に関わるものである。
電子ビーム溶解法によって鋳造した、Ｏ成分が０．０３８mass％、Ｆｅ成分が０．０３５mass％で、スラブ厚さ６３０mmの範囲にある矩形断面チタンスラブを、β域（９５０℃）に加熱し、圧下比が４．２の分塊圧延を行い、その後直ちにβ域加熱終了温度〜７００℃の温度範囲を冷却速度３５０℃／ｈｒで冷却した。このようにして製造した板厚１５０mmの分塊圧延スラブを複数枚準備して、その後α域（８４０℃）に加熱し、分塊スラブの板幅方向（Ｃ方向）に圧下比２．２の粗熱延延を行い、引き続いて仕上げ熱延として、粗熱延との圧延方向が直交する（Ｌ方向）クロス熱延を、温度範囲が６７０℃〜７２０℃の間においてクロス圧延比がＣ／Ｌ＝５／１０となるように行った。さらに真空クリープ板形状矯正ＶＣＦ焼鈍を、温度範囲が６１０℃、保定時間が５ｈｒとなる条件で行った。
【００６９】
これらのＶＣＦ焼鈍材に対して、板長手方向に圧延率２０〜６０％の範囲の冷間圧延を行い、５８０〜７５０℃の温度範囲で焼鈍を行い、得られた焼鈍板の表面下１／３板厚層板面におけるマクロ組織および平均結晶粒径を評価した。その結果を表７に示す。
本発明例のＥ６〜Ｅ８はマクロ模様もなく、平均結晶粒径も４０μｍ以下となり、電解銅箔ドラム用チタン厚板として良好な素材となる。
【００７０】
これに対して比較例のＥ１〜Ｅ４は、焼鈍後の冷間圧延率が本発明範囲の下限２５％に満たないため、マクロ模様が表れた。比較例Ｅ５，Ｅ１１は、焼鈍温度が本発明範囲より低いため、マクロ模様が表れた。比較例Ｅ９，Ｅ１０は、焼鈍温度が本発明範囲より高いため、結晶粒が粗大化している。保定時間が３５ｍｉｎである比較例Ｅ１２〜Ｅ１４は、表面酸化が過剰で宜しくない。また冷間圧延率が６０％と高い比較例Ｅ１５〜Ｅ１８は、板形状が不良であった。
【００７１】
【表７】

【００７２】
【発明の効果】
本発明は以上のような構成により、熱間圧延において設備上の制約あるいは製造上の難点を有する、いわゆる二段加熱圧延法によらない一段加熱圧延法によって、板表層部位から１／３板厚層部位にわたる金属組織が均一緻密であり、マクロ模様（スクラッチ疵）の原因となるマクロ的不均一組織のない、電解銅箔製造のための電着ドラム用チタン厚板の製造が可能となる。
【図面の簡単な説明】
【図１】α域一段加熱連続一方向圧延材のミクロ組織を示す図である（観察倍率１００倍）。
【図２】α域一段加熱連続クロス圧延材のミクロ組織を示す図である（観察倍率１００倍）。
【図３】α域二段加熱クロス圧延材のミクロ組織を示す図である（観察倍率１５倍）。
【図４】α域二段加熱クロス圧延材のミクロ組織を示す図である（観察倍率１００倍）。
【図５】ＶＣＦ焼鈍温度、保定時間とＶＣＦ焼鈍板のミクロ組織平均結晶粒径の関係を示す図である。
【図６】最終焼鈍温度、保定時間と最終焼鈍板のミクロ組織平均結晶粒径の関係を示す図である。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to titanium for a copper foil production drum having an excellent surface layer structure and a method for producing the same.
[0002]
[Prior art]
A printed wiring board used by being incorporated in an electronic component is made by bonding a copper foil and an insulating substrate, printing a conductor pattern on the surface, and removing unnecessary portions by etching. Electrolytic copper foil is mainly used for this printed wiring board, but such electrolytic copper foil is a rotating high-quality copper raw material dissolved in sulfuric acid solution, with lead or other insoluble metal as the anode and facing it. A copper foil is electrochemically electrodeposited on the drum using the drum as a cathode, and this is continuously peeled off from the rotating drum and wound into a roll. As a material for the cathode, titanium has been frequently used in recent years due to corrosion resistance and peelability of the electrodeposited metal.
[0003]
When copper foil is used for printed wiring, the surface roughness of the copper foil is formed by etching. Wiring This is an important factor that affects the accuracy of the pattern (line width: 0.1 to 0.5 mm). This copper foil surface roughness directly inherits the polished and leveled roughness of the titanium cathode drum on which the copper foil is electrodeposited.
[0004]
In addition, the polished and smoothed cathode drum may be gradually corroded in the electrolytic solution, and may be in a state where it is subjected to etching used in metal structure inspection. In such a case, there is a problem that the metal structure macro pattern on the drum surface is transferred onto the copper foil, the uniformity of the finished surface roughness is impaired, and the printed wiring is poorly etched. The transferred pattern of the metal structure is called a scratch wrinkle, shirakmo, swordfish or the like because of its form (hereinafter collectively referred to as a scratch wrinkle).
Furthermore, since the surface of the drum becomes rough due to electric sparks during use, the drum surface is polished and leveled many times. As a result, the drum surface is ground gradually and the new surface becomes the drum surface. Accordingly, the drum material is required to have uniformity in the thickness direction in addition to uniformity on the plate surface.
[0005]
The cause of scratch defects is empirically known to be due to the so-called macroscopic non-uniform structure in the drum material, and the raw material of the drum manufacturing material, that is, slabs and hot-rolled rough pieces, etc. Attempts have been made to homogenize the macrostructure present in the. For example, investigations on "working heat treatment methods" such as forging conditions for ingots or hot strip rolling, high temperature heat treatment of split slabs, high temperature long time annealing of hot rolled thick plates, etc. have been conducted, and appropriate improvements can be obtained. Has been. However, the quality requirements for the copper foil are becoming higher and the density of the metal structure is no longer satisfied by the existing methods.
[0006]
Conventionally, in a plate material generally called a titanium expanded material, mechanical properties such as tensile strength and elongation and crystal grain size defined in JIS standards and the like have been required as material performance characteristics. However, like the drum material for copper foil production, it is required to avoid the remaining macroscopic non-uniform structure inherited from the raw material (ingot) in the processing process, or the macro structure is uniform in the surface layer region. There was no case where the density was required. In other words, even if the macro structure in the surface layer portion is non-uniform and lacks in density, there is no particular problem in applications other than the copper foil manufacturing drum as long as the above standards are satisfied.
[0007]
On the other hand, in the copper foil manufacturing process, a titanium drum rotates in a copper sulfate aqueous solution, and copper is electrodeposited on the surface thereof. At that time, the drum surface is corroded in the copper sulfate solution, and preferentially erodes the crystal grain boundaries and the second phase interface having a high lattice defect density, and the boundary portions of the adjacent orientation group grains called colonies. In this way, the eroded macroscopic non-uniform pattern is transferred to the surface of the copper foil. In the colony in this case, the shape, size, and direction of the etch pit facets generated by the corrosive action are aligned, and when they are transferred to the copper foil, there are various kinds from a very slight level to a clear level. It is thought that this causes a dull quality of the copper foil called the above-mentioned scratched wrinkle.
[0008]
As a conventional example, in Japanese Patent Application Laid-Open No. 60-9866, in order to eliminate scratches, the heating temperature in the forging and rough hot rolling is set to 950 ° C. or more, and the rolled sheet is again used during the finishing hot rolling. A uniform fine macro is formed by a “β region-α region two-stage heating cross rolling method” in which a cross hot rolling is performed in which the hot rolling is performed in a direction orthogonal to the rolling direction at the time of rough hot rolling after heating to 700 ° C. or less. A method for producing patterned titanium and titanium alloy plates is disclosed.
According to this, the cause of the non-uniform macro pattern (hereinafter referred to as macro defect) is (a) macro defect due to residual cast structure, (b) macro defect due to residual transformation structure, (c) macro defect due to residual coarse crystal grains, (D) It is classified into four types of band-like macro defects parallel to the rolling direction caused by rolling, and the above-mentioned “β region-α region two-stage heating cross rolling method” includes (a), (b), ( This is an effective method for eliminating all of c) and (d).
[0009]
However, in this method, it is necessary to insert the rolled sheet into the heating furnace for finishing hot rolling again during the rolling after rough rolling, which means that two or more heating furnaces having different heating temperatures should be installed. Since this is a premise, it imposes significant restrictions on the production equipment mill configuration. In addition, this method is difficult to control the plate shape because the rolling deformation resistance is high because the finish hot rolling is performed in a low temperature range of 700 ° C. or less, and the rolling is performed in a low temperature range below the recrystallization temperature. In many cases, the above-mentioned type (d) macro defects tend to occur more frequently.
[0010]
On the other hand, in Japanese Patent Application Laid-Open No. 9-176809, (1) when performing slabs obtained by rough-rolling and finish hot-rolling by using the “α region one-stage heating cross rolling method”, that is, by forging or by rolling. The slab heating temperature during rough hot rolling is set to a temperature range of 700 ° C. or more to less than the β transformation temperature, rolling is performed in the longitudinal direction and the width direction orthogonal to each other, and the amount of processing true strain due to rolling in the longitudinal direction is ε _L , The true strain amount due to rolling in the width direction is ε _w When ε between both processing true strain amount _L × ε _w A method for producing a titanium or titanium alloy plate without a non-uniform macro pattern is disclosed, wherein the rolling is performed under a condition that satisfies ≧ 0.25.
Furthermore, (2) as a slab to be subjected to rough hot rolling, use a slab whose crystal grains are refined by heating to a temperature higher than the β transformation temperature after being hot-worked below the β transformation temperature, or 3) A manufacturing method is disclosed in which a hot rolled sheet is annealed at 800 ° C. or lower and then cold rolled and annealed to add an average crystal grain size of 50 μm or less.
[0011]
However, among these, (1) defines the total rolling reduction in the sheet longitudinal direction and sheet width direction in cross rolling, but does not define an appropriate rolling temperature range in which cross rolling should be performed. Is extremely insufficient. In (2), when the slab is reheated in the β single phase region, grain growth occurs remarkably, and the refinement of the structure is not achieved. Further, in (3), details of the temperature at which the hot-rolled sheet is annealed, the rolling ratio of the subsequent cold rolling and the temperature of the final annealing, and the holding time are not specified, and a titanium sheet having an average particle diameter of 50 μm is obtained. Insufficient technical requirements are defined.
[0012]
[Problems to be solved by the invention]
In view of the above circumstances, an object of the present invention is to provide a titanium foil for a copper foil production drum excellent in a surface layer structure, which does not exhibit a macroscopic non-uniform structure that causes scratches after corrosion in the surface layer part of the plate, and its A manufacturing method is provided.
In particular, with respect to the production method, a so-called single-stage hot rolling method that does not rely on the so-called two-stage hot rolling method, which has the above-described equipment limitations or manufacturing difficulties in hot rolling, and an integrated production method including the same are provided. Is.
[0013]
[Means for Solving the Problems]
As a result of intensive studies on the relationship between the surface layer structure of titanium used for the copper foil manufacturing drum and scratches, etc., the inventors have found the conditions to be satisfied by the surface layer structure, and the components for this and restrictions on equipment The present invention has been completed by limiting the production method with a small amount, and the gist thereof is as follows.
[0014]
(1) Mass%,
Fe: 0.04% or less, O: 0.1% or less
In the surface layer portion parallel to the plate surface extending from the outermost surface to 1/3 plate thickness, the average particle size of the α phase is 40 μm or less, and the α phase texture is [ 0001] 0 ° to ± 45 ° TD and [0001] 0 ° to ± 25 ° RD (where TD is the plate width direction and RD is the rolling direction) Titanium for manufacturing drums.
[0015]
(2) When producing titanium for a copper foil production drum as described in (1) above, a rectangular cross-section slab having a thickness of 300 mm or more melted and cast by an electron beam melting method is heated in the β region, and the reduction ratio in the β region. Immediately after forming a β-phase recrystallized structure by performing 3 or more ingot rolling or ingot forging, immediately cool the range of the β region processing end temperature to 700 ° C at a cooling rate of 200 ° C / hr or more. Then, it is heated to 880 ° C. or less and subjected to rough hot rolling, and the cross hot rolling is performed by rolling in the direction perpendicular to the rolling direction of the rough hot rolling without reheating after the rough hot rolling. Finishing hot rolling so as to be 10 to 6/10 is performed in a temperature range of 650 to 750 ° C., and further annealing is performed in a temperature range of 570 to 680 ° C. for 1 to 10 hours. A method for producing titanium for a copper foil production drum having an excellent surface layer structure.
( 3 ) After the annealing, the rolling rate is further cold-rolled at 25 to 50%, and then annealing is performed at a temperature range of 600 to 710 ° C. for 10 to 30 minutes. 2 The manufacturing method of titanium for copper foil manufacturing drums excellent in the surface layer part structure described in).
[0016]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, the present invention will be described in detail.
As a result of intensive studies to elucidate the mechanism of scratch crease formation, particularly its metallurgical cause, the inventors first found that the structure of a macroscopically heterogeneous structure in a hot-rolled sheet that directly causes scratch crease. The following important findings were obtained regarding the relationship between science and crystallography. That is, the inventors conducted detailed research on hot rolling texture formation in thick titanium rolling and the like.
[0017]
as a result,
(1) A slab heating temperature is β region, and β region heating β to α region continuous rolled material which is continuously rolled in one direction (or reverse direction) parallel to the slab longitudinal direction in the β to α region, [0001] shows a so-called transverse-texture in which the axis is inclined by 90 ° in the plate width direction. On the contrary,
(2) When the heating temperature of the slab is set to the α region and rolling is continuously performed in one direction (or the reverse direction) in the α region, the [0001] axis is inclined by 35 to 45 ° in the plate width direction. Indicates TD-texture. further,
(3) When cross rolling is performed by rotating the rolling direction by 90 ° in a low temperature range (about 600 to about 750 ° C) in the α range, the [0001] axis is almost parallel to the plate normal direction (within about ± 25 °) Form the so-called Center-Pole-texture.
It revealed that.
Titanium plates produced by the method of (3) have a high tendency to form a macroscopic non-uniform structure in the cases of (1) and (2). Found that the fine and macroscopic heterogeneous structure was reduced to a very slight degree.
[0018]
The metallurgical reason why the macroscopic heterogeneous structure is reduced in the titanium plate produced by the method (3) is considered as follows.
That is, generally, when unidirectional (or reverse) rolling is performed, a shear slip deformation band having an inclination angle of about 45 ° with respect to the rolling surface is formed, but when cross rolling is performed, the rolling surface normal axis is further increased by 90. A rotated 45 ° shear slip deformation zone is additionally introduced. Since rolling twins are formed in rolling in the low temperature region of the α region, paying attention to a single α-phase crystal grain, the structure of the crystal grain is extremely high due to cross-slip and twin deformation caused by cross rolling. Finely divided. The processing strain based on such slip deformation increases the recrystallization driving force, promotes recrystallization during rolling, and forms a uniform fine grain structure. In addition, twins having a high dislocation density also form pseudo grain boundaries, which also promotes refinement of the structure.
[0019]
Thus, if cross rolling is performed in the low temperature region of the α region, the breakage and refinement of the macroscopic nonuniform structure is promoted, and this is an extremely effective method for reducing scratches. In other words, in order to eliminate the macroscopic non-uniform structure of the titanium plate, it is shown that rolling to form a center-pole-texture in the hot finish rolling stage is a very important requirement. .
[0020]
As described above, in order to eliminate the macroscopic non-uniform structure, it is extremely important to perform cross-rolling from rough hot rolling to finish hot rolling, and (c) not to cause residual coarse crystal grains. In order to more effectively express the slab structure, the slab structure before rough hot rolling shows (a) a macro defect due to residual cast structure and (b) a coarse non-uniform structure that causes macro defect due to residual transformation structure. must not.
[0021]
The inventors have studied in detail the relationship between the shape of the slab to be subjected to rough hot rolling, the slab lump pressure condition, the slab structure, etc., and the formation of a macroscopic heterogeneous structure in the hot-rolled sheet, and the electron beam melting. Thickened rectangular cross-section titanium slab with a slab thickness of 300mm or more melted and cast by the method is heated in the β region and subjected to ingot rolling or ingot forging at a reduction ratio of 3 or more in the β region to form a β phase recrystallized structure Immediately after that, when the range of the β region processing end temperature to 700 ° C. is cooled at a cooling rate of 200 ° C./hr or more to form a fine transformation structure based on the β → α transformation, It was found that the heterogeneous structure decreased.
[0022]
The ingot produced by the consumable electrode arc melting (VAR) method usually has a cylindrical shape with a cross-sectional diameter of about 500 to 800 mmφ. Therefore, in order to hot-roll this cylindrical ingot, Or it is necessary to process into a rectangular section slab by partial forging. In that case, compression processing is performed by rolling or forging from the direction perpendicular to the cylindrical axis direction of the cylindrical ingot (normal direction of the outer peripheral surface). At this time, processing corresponding to the end of the slab cross-sectional width direction after processing In the vicinity of the outer peripheral portion of the front cylindrical ingot cross section, it is difficult to ensure a sufficient reduction ratio as compared with the vicinity of the ingot center portion. As a result, there is a difficulty that the cross-sectional structure of the slab after the partial rolling / forging shows a non-uniform structure as a whole.
[0023]
In the VAR ingot, ingot rolling and forging are generally heated in the β range and processed from the β range to the α range. This is because when the α region heating α region processing is performed, there is no opportunity for the material to undergo a structure refining action based on β → α transformation, and as a result, (a) occurrence of macro defects due to residual cast structure cannot be avoided. As described above, the outer periphery of the VAR cylindrical ingot is required to be processed with a low reduction ratio, and a coarse β-phase non-recrystallized structure is likely to be present at the same part. (B) Macro defects due to residual transformation structure also occur. The tendency to do is great.
[0024]
On the other hand, in the present invention, a rectangular cross-section slab to be cast into a rectangular slab mold after melting by hearth is manufactured by an electron beam melting (EBR) method. There is no problem that the subsequent slab cross-sectional structure becomes non-uniform.
[0025]
In addition, since the plate thickness in the ingot rectangular cross section is 300 mm or more, desirably about 600 mm thick, the rolling finish plate thickness of the ingot rolled slab is assumed to be 100 mm, in β zone rolling after β zone heating. It is possible to secure a reduction ratio from 3 (processing rate 66%) or more to a maximum of 6 (processing rate 83%). Further, since the thickness of the ingot is sufficiently thick, rolling in the complete β region is performed without the temperature during rolling falling to the α region.
As described above, when processing with a reduction ratio of 3 or more in the β region, the β phase is recrystallized during processing and becomes finer, and further transformed β structure by β → α transformation in the subsequent cooling process. (Fine α lamellar structure) is obtained. As a result, sufficient structure refinement and randomization of crystal orientation are realized as a slab structure used in the hot rolling process.
[0026]
In addition, in the slab cooling process, if staying in a temperature range from 700 ° C. to a temperature range from 700 ° C. at which the α-region grain growth is remarkable, the α lamellar structure becomes coarse and undesirable. The correct temperature range needs to be cooled as quickly as possible.
The inventors have found that a fine transformation structure based on the β → α transformation can be formed by cooling at a cooling rate of 200 ° C./hr or more, particularly in the range of the β-region working end temperature to 700 ° C. On the other hand, when the slow cooling is performed at a cooling rate of less than 200 ° C./hr, the transformation structure is coarsened to cause macro defects.
[0027]
Furthermore, in the present invention, in the hot rolling of a titanium slab cast, split rolled or forged by the above-mentioned method, the heating temperature in the rough hot rolling is set to 880 ° C. or less, and reheated after the rough hot rolling. It is preferable to adopt a production method in which finish hot rolling is performed continuously at 650 to 750 ° C.
[0028]
In this finish hot rolling, a cross hot rolling that is rolled in a direction perpendicular to the rolling direction of the rough hot rolling is taken in and rolled so that the cross rolling ratio becomes 1/10 to 6/10. Cross rolling ratio = previous stage total reduction ratio before cross rolling / rear stage total reduction ratio after crossing.
In this case, the reason for setting the slab heating temperature in rough hot rolling to 880 ° C. or lower is that when heated to 880 ° C. or higher corresponding to the β transformation temperature, the fine α lamellar structure is transformed into β phase and β phase grain growth occurs. This is because it is necessary to avoid this.
[0029]
The inventors manufactured a slab (150 mm (t) x 1260 mm (w) x 850 mm (L)) produced by EBR melted ingot having a thickness of 600 mm by β-zone heating and ingot rolling and quenching after finishing the β zone. Ten pieces of rectangular parallelepiped rolling blocks having dimensions of 150 mm (t) × 170 mm (w) × 250 mm (L) are collected from the β block rolling reduction ratio: 4.0), and these rolling blocks are collected in the α region ( 850 ° C.) and the following three types of rolling experiments were conducted to examine detailed rolling conditions.
[0030]
(A): α region one-stage heating continuous unidirectional rolling (second-stage rolling temperature control)
In the α region one-stage heating continuous unidirectional rolling, the former rough rolling is rolled in 150 passes from 30 mm to 30 mm in 7 passes, and then the finishing finish rolling start temperature is changed in the range of 500 to 750 ° C. Four-pass rolling was performed in the direction (L direction) from 30 mm thickness to 9.0 mm thickness. In this case, the former stage rolling has a rolling rate of 80% (a reduction ratio of 5.0), and the latter stage rolling has a rolling ratio of 70% (a reduction ratio of 3.3). The total rolling rate from the start of rolling to the end of rolling is 94% (rolling ratio 16.7). FIG. 1 (observation magnification × 100) is a result of microstructural observation of a 1 mm thick layer portion parallel to the plate surface below the surface of the rolled material.
[0031]
In the case of the high-temperature rolled material shown in (1) and (2) of FIG. 1 where the rolling start temperature of the latter stage rolling is 750 ° C. and 700 ° C., fine recrystallized grains are generated, but the latter stage rolling temperature is As (3) 600 ° C. and (4) 500 ° C., the non-recrystallized mixed grain structure including rolled twins is shown. Also in the macro etch structure observation, it was confirmed that macro grains stretched in the direction parallel to the rolling direction (L direction) were particularly generated in the low-temperature post-rolled rolled material (600 ° C., 500 ° C.). Therefore, from the microstructural observation and the macrostructural observation results, it is easy to form a remarkable macro-uniform structure that is parallel to the rolling direction in the α-zone one-stage heating continuous unidirectional rolling, and this tendency becomes conspicuous with a decrease in the subsequent rolling start temperature. It is clear that
[0032]
(B): α region one-stage heating continuous cross rolling (second-stage cross rolling temperature control)
In the α-zone one-stage heating continuous cross rolling, the preceding rough rolling is rolled in the C direction in four passes from 150 mm thickness to 84 mm thickness in the 170 mm (W) direction, and then rotated immediately by 90 degrees without cooling and reheating. Then, the latter-stage rolling start temperature was changed in the range of 500 to 750 ° C., and 8-pass cross rolling was performed in the L direction from 84 mm thickness to 9.0 mm thickness. In this case, the first-stage rough rolling has a rolling rate of 44% (rolling ratio 1.78), and the latter-stage finish rolling has a rolling ratio of 89.2% (rolling ratio 9.33). Therefore, the cross ratio C / L = 1 / 5.24, and the total rolling rate from the start of rolling to the end of rolling is 94% (rolling ratio 16.7).
[0033]
FIG. 2 (observation magnification × 100) is a result of microstructural observation of a 1 mm thick layer portion parallel to the plate surface below the surface of the rolled material. As can be seen in FIG. 2, (1) EB5, post-rolling start temperature 750 ° C., (2) EB6, post-rolling start temperature 700 ° C., all of the materials show a fine recrystallized equiaxed structure and are extremely good. Become an organization. (3) EB7, post-rolling start temperature 600 ° C. material, non-recrystallized stretched grains slightly parallel to the rolling direction are observed.
On the other hand, also in the macroetch structure observation, the EB5 and EB6 materials showed a very uniform macrostructure without showing a nonuniform macrostructure containing unrecrystallized α-stretched grains. On the other hand, the macro structure of the EB7 material exhibited a macro-inhomogeneous structure including a few unrecrystallized α-stretched grains. As described above, in the α region one-stage heating continuous cross rolling, when the post-rolling start temperature is set to 700 to 750 ° C., a very uniform macrostructure is obtained, and a remarkable improvement effect is obtained as compared with the α region one-stage heating continuous unidirectional rolling. It can be seen.
[0034]
(C): α region two-stage heating cross rolling (second-stage cross rolling temperature control)
In the α-region two-stage heated cross rolling, the preceding stage rough rolling was rolled in the C direction in 5 to 7 passes from 150 mm to 60 mm (EB9) or 35 mm (EB10) and water cooled. Thereafter, the rolled plate was cut, reheated to 700 ° C., and subjected to L total cross rolling in which the rolling direction was rotated 90 degrees from the rolling direction of the preceding rough rolling. At that time, the rolling start temperature of the latter finish rolling was fixed at 700 ° C., and cross rolling of 5 to 7 passes was performed in the L direction from 60 mm thickness or 35 mm thickness to 9.0 mm thickness.
In this case, the former rolling is a rolling rate of 60% (rolling ratio 2.5) or rolling rate 77% (rolling ratio 4.28), and the latter rolling is a rolling rate of 89.2% (rolling ratio 9.33) or rolling rate 74. 0.2% (reduction ratio 3.8). Therefore, the cross ratio of EB9 is C / L = 1 / 2.66, or EB10 is C / L = 1.0 / 0.90. The total rolling rate from the start of rolling to the end of rolling is 94% (rolling ratio 16.7).
[0035]
FIG. 3 (observation magnification x15) and FIG. 4 (observation magnification x100) are the results of microstructural observation of the 1 mm thick layer portion parallel to the plate surface below the surface of the rolled material. In the figure, (1) In EB9, C / L = 1.0 / 2.66, macroscopic stretched grains are almost eliminated, and a fine grain uniform structure is formed. Continuous crossing at 700 ° C. without reheating as shown in FIG. Rolling EB6 (cross ratio C / L = 1.0 / 5.24) performed and EB9 (cross ratio C / L = 1.0 / 2.66) subjected to 700 ° C. reheat cross-rolling shown in FIG. When comparing the two, they showed almost good fine uniform equiaxed structure, and it was confirmed that performing the cross rolling in this temperature range is extremely effective for eliminating the macroscopic nonuniform structure.
[0036]
On the other hand, the EB10 material shown in FIGS. 3 (2) and 4 (2), in which the cross ratio was reinforced and C / L = 1.0 / 0.90, formed α-stretched grains, Exhibits a heterogeneous structure. The reason for this is that the EB10 material has a condition where the cross ratio is close to 1.0 / 1.0, but since the rolling reduction at the latter stage is as low as about 3.8, it is difficult to eliminate the α-drawn grains. It seems that it was enough. On the other hand, in the macroetch structure observation in which the observation area was further expanded, the EB9 and EB10 materials showed a non-uniform macrostructure including unrecrystallized α-stretched grains in some places.
[0037]
From the results of (A), (B), and (C) described above, the rectangular cross-section slab cast by electron beam melting, as described above, is subjected to partial rolling (or forging) in the β region and cooled to the titanium slab, In the step of sequentially performing the rough hot rolling and the finishing hot rolling, the heating temperature in the rough hot rolling is set to 880 ° C. or less, and in the finishing hot rolling continuously performed without reheating after the rough hot rolling, The cross hot rolling with the rolling direction orthogonal to each other has a temperature range of 650 ° C. to 750 ° C., and the cross rolling ratio (= the total pre-rolling ratio before the cross rolling / the total post-crossing post-crossing reduction ratio) is 2/10 (about 1 0.0 / 5.24) to 4/10 (about 1.0 / 2.66) has been found to be most appropriate. In this case, even when the upper limit of the cross rolling ratio was expanded to 6/10 and the lower limit was expanded to 1/10, the macroscopic nonuniform structure could be eliminated.
[0038]
Since the titanium thick plate (as hot-rolled material) manufactured by the above method often has a poor shape such as flatness, it is annealed in a temperature range of 570 to 680 ° C. For annealing, it is preferable to perform an annealing treatment that also serves as shape correction using a vacuum creep straightening (VCF) equipment.
Table 1 shows the relationship between the as-rolled material, various VCF annealing temperatures and the macrostructure, and FIG. 5 shows the relationship between the VCF annealing temperature and the average crystal grain size in the microstructure of the VCF annealed material.
[0039]
[Table 1]

[0040]
As is apparent from Table 1, macroscopic nonuniformity remained in the as-rolled material, but when annealing was performed at a temperature of 570 ° C. or higher, recrystallization was completed and the macroscopic nonuniform structure disappeared completely. The holding time at the holding temperature of the VCF annealing is in the range of 1 to 10 hours, but when the average crystal grain size of the VCF annealed material becomes coarser than 40 μm as shown in FIG. 5, the VCF annealed material is used as a drum material. In this case, the drum polishability becomes poor.
[0041]
In addition, when manufacturing a titanium thick plate for a copper foil drum by adding a cold rolling process and an annealing process following VCF annealing, the average crystal grain size of the material before cold rolling is excessively coarse. Since it causes a problem such as causing a macro-inhomogeneous structure, it is appropriate that the average particle diameter of the material is substantially in the range of 10 to 70 μm while being annealed by VCF.
[0042]
As shown in FIG. 5, the VCF annealing temperature range is set to 570 to 680 ° C., and the holding time range is set to 1 to 10 hours, so that cold rolling can be performed on the material as it is after VCF annealing as described later. It is possible to provide a titanium thick plate for electrodeposition drums for the production of electrolytic copper foil without fine, uniform and macro patterns even when manufacturing titanium thick plates for copper foil drums by adding processes and annealing processes. To. In addition, also in the case of normal annealing which is not VCF, temperature and time are performed according to the above conditions.
[0043]
The titanium thick rolled sheet subjected to the α region one-stage heating and continuous cross rolling is subjected to VCF annealing for 5 hours holding at 630 ° C., and after cold rolling with a total rolling rate of about 40%, 600 to 750 ° C. Annealing for holding for 10 to 30 minutes in the temperature range was added. FIG. 6 is a diagram showing the relationship between the average crystal grain size of the α phase, the annealing temperature, and the annealing holding time observed from the cross-sectional metal microstructure in the 1/3 sheet thickness phase region from the plate thickness outermost layer region of such an annealed plate. is there.
[0044]
In FIG. 6, the area surrounded by the hatched area is the appropriate temperature and holding condition range for the final annealing in the present invention. For example, when the average crystal grain size of the final annealed titanium plate is 5 μm or more, the macro non-uniform structure is almost eliminated. When the average crystal grain size is more than 40 μm, the macroscopic non-uniform structure is eliminated, but the coarse crystal grain size excessively increases the roughness of the copper foil surface, resulting in poor drum polishability. This is not desirable. Accordingly, the average crystal grain size may be in the range of 5 to 40 μm. However, the retention time for realizing this is, for example, 10 minutes or less, and a uniform temperature distribution in the plate surface cannot be achieved, and the crystal grain size Is not good because it mixes.
[0045]
Further, if the holding time is 30 minutes or more, the surface oxidation becomes excessive, the load of the subsequent deoxidation film process becomes large, and the productivity is not good.
For the above reasons, the final annealing condition after cold rolling was limited to 10 to 30 min in the temperature range of 600 to 710 ° C. Further, regarding the cold rolling rate before the final annealing, it is desirable that the total rolling rate is 25 to 50% in the long side direction of the thick plate on which the VCF annealing is performed. If the rolling rate is less than 25%, the metal structure of the final annealed plate is likely to be mixed and a macro non-uniform structure may remain, which is not desirable. If the rolling rate is 50% or more, the rolled plate shape becomes poor, which is also undesirable.
[0046]
Next, since the titanium thick plate manufactured by the series of manufacturing methods based on the present invention as described above is subjected to cross hot rolling based on an appropriate cross rolling ratio in an appropriate temperature range, rolling or annealing is performed. The texture obtained in the plate is characterized in that, for example, Center-Pole-texture, Split-TD-texture, Split-RD-texture, and Transverse-texture are mixed in the hot-rolled plate.
[0047]
For this reason, Split-TD-texture (+ Split-RD-texture) is likely to be formed in the VCF annealing or in the VCF annealing + cold rolling + final annealing. Therefore, regardless of the thermomechanical process applied to the titanium material, the texture from the outermost layer part of the plate thickness to the one-third plate thickness layer part is determined by Center-Pole-Texture, Split-TD-Texture. And Split-RD-texture single or composite texture, that is, [0001] 0 ° to ± 45 ° TD and [0001] 0 ° to ± 25 ° RD (where TD is the sheet width direction and RD is the rolling direction) ) Has a texture that is the main orientation, and the average crystal grain size of the surface layer α-phase grains in the surface to 1/3 plate thickness portion is 40 μm or less, and the fine grained uniform macroscopic non-uniform electrolysis Any titanium plate for copper foil drums may be used.
[0048]
Therefore, as described above, the method described in detail in the embodiment of the present invention is most suitable for the manufacture thereof, but other methods are not particularly limited. The chemical composition of pure titanium to which the present invention is applied is desirably O ≦ 0.1 mass% and Fe ≦ 0.04 mass%.
[0049]
The reason for O ≦ 0.1 mass% is that when a thick titanium plate is formed into a drum-shaped ring and both ends are welded together, about 200 ppm of oxygen is additionally mixed during welding (oxygen pickup). Therefore, the welded part may have a hardness higher than that of the base metal part, and if O is more than 0.1 mass%, the hardness increase of the welded part becomes more prominent, and the ring-welded welded part is cold worked and annealed. Even if the surface is polished and polished, a convex surface defect having a linear height of about 0.5 μm remains along the weld line, and the concave linear defect is transferred to the copper foil. Because.
[0050]
Further, the reason for Fe ≦ 0.04 mass% is that when Fe is higher than 0.04 mass%, pitting corrosion-like pitting based on Fe that is unevenly distributed in crystal grain boundaries becomes remarkable, and this is transferred to the copper foil. Therefore, convex surface defects are formed on the copper foil.
[0051]
【Example】
Example 1
This embodiment relates to claim 1 of the present invention.
20 types of high-purity titanium EBR titanium ingots or VAR ingots with O of 0.02 to 0.13 mass% and Fe of 0.013 to 0.06 mass% are manufactured by dissolution. Table 2 in Table 2 was subjected to the heat treatment shown in (1) to (6) to produce a titanium thick plate.
[0052]
For example, in the manufacturing process (1), an EBR ingot (dimension: 600 mm thickness × 1300 mm width × 2000 mm length) is heated to the β region (1000 ° C.), and the rolling at a reduction ratio of 4.0 is performed to a thickness of 150 mm in the β region. went.
At this time, the cooling rate between the β region processing end temperature (900 ° C.) and 700 ° C. was set to 250 ° C./hr. After that, it is heated to the α region (850 ° C.), subjected to rough hot rolling with a reduction ratio of 3.0 in the plate width direction (C direction) of the slab, and then as the finish hot rolling, with respect to the direction of rough hot rolling. Cross hot rolling in which the rolling directions are orthogonal (L direction) was performed so that the cross rolling ratio was C / L = 6/10 in the temperature range of 700 to 750 ° C. Thereafter, vacuum creep plate shape correction VCF annealing was performed (650 ° C. × 1 hr).
[0053]
In the manufacturing process (2), the EBR ingot (dimension: 600 mm thickness x 1300 mm width x 2000 mm length) is heated to the β region (1000 ° C), and ingot is rolled at a reduction ratio of 4.0 to a thickness of 150 mm in the β region. It was. At this time, the cooling rate between the β region processing end temperature (900 ° C.) to 700 ° C. was set to 300 ° C./hr.
After that, it is heated to the α region (850 ° C.), subjected to rough hot rolling with a reduction ratio of 2.0 in the plate width direction (C direction) of the slab, and subsequently the rolling direction with the rough hot rolling as the finish hot rolling is The cross (L direction) cross hot rolling was performed so that the cross rolling ratio was C / L = 4/10 in the temperature range of 700 to 750 ° C. Further, vacuum creep plate shape correction VCF annealing was performed (650 ° C. × 1 hr). Thereafter, cold rolling at a rolling rate of 40% was performed, and final annealing (660 ° C. × 20 min) was added. Hereinafter, manufacturing processes (3) to (6) are the methods described in Table 2.
[0054]
[Table 2]

[0055]
The plate surface of the obtained titanium thick plate was polished to a 1 mm layer plate surface and a 1/3 plate layer surface under the surface, respectively, and corroded with nitric hydrofluoric acid, and macroscopic visual inspection was performed. Table 3 shows the evaluation results. In Table 3, the crystal grain size means the average crystal grain size of the 1/3 plate thickness layer part plate surface, and the texture is the main orientation (maximum) obtained by measuring the (0001) positive pole figure of the plate surface layer of the same part. Intensity extreme density peak position) and sub-direction (polar density peak positions of second and third intensities).
Here, Split-TD (± 42 °) is Split-TD-texture whose [0001] axis is inclined by ± 42 ° in the TD direction, Center-Pole is Split-RD (0 ° to ± 25 °), and Split-TD -Center (Pole-Texture) in which the [0001] axis where TD (0 ° to ± 25 °) is mixed is within 25 ° from the ND (plate normal) direction. Transverse means Transverse-texture in which the [0001] axis is inclined by 70 ° or more in the TD direction.
[0056]
The titanium plates of the present invention of sample symbols A1 to A9 manufactured in the manufacturing steps (1) and (2) shown in Table 3 show almost no macro pattern, and the crystal grain size is 40 μm or less and the α phase texture. [0001] 0 ° to ± 45 ° TD and [0001] 0 ° to ± 25 ° RD. Sample symbols A10 to A20 of the comparative example are examples in which a mild and sometimes severe macro pattern is generated, and neither the crystal grain size nor the α phase texture satisfies the above requirements.
[0057]
[Table 3]

[0058]
(Example 2)
This embodiment relates to claim 2 of the present invention.
A rectangular cross-section titanium slab cast by an electron beam melting method with an O component of 0.04 mass%, an Fe component of 0.03 mass%, and a slab thickness in the range of 250 to 650 mm is obtained in an α region (850 ° C.) and β It heated to the area | region (950 degreeC, 1000 degreeC), and the partial rolling was performed in the range whose rolling ratio is 1.5-5.0. Thereafter, the temperature range from 700 ° C. to β-region ending temperature was immediately cooled at a cooling rate of 100 to 500 ° C./hr.
These various block-rolled slabs were subjected to thick plate rolling and VCF annealing in the same manner as in “Process (1)” in Example 1, and the obtained titanium thick plate surface was reduced to 1/3 plate. The thick layer portion was polished and corroded with nitric hydrofluoric acid, and the macro pattern was visually inspected.
[0059]
In addition, a cylindrical ingot having a diameter of 800 mm (0.04 mass% O, 0.031 mass% Fe) in which O and Fe components were subjected to VAR melting with substantially the same amount was heated to the β region (1000 ° C.) and reduced. The ratio rolling (diameter portion reduction ratio) is performed in the range of 2.8 to 4.5, and immediately thereafter, the temperature range of the processing end temperature in the 700 ° C. to β range or α range is set to the cooling rate: 100 to 300 ° C. Cooled at / hr. These various block-rolled slabs were subjected to thick plate rolling and VCF annealing in the same manner as in “Process (1)” in Example 1, and the plate surface of the obtained titanium thick plate was reduced to 1/3 plate thickness. The layer portion was polished and corroded with nitric hydrofluoric acid, and macroscopic visual inspection was performed. Table 4 summarizes the manufacturing conditions and macrostructure evaluation results of these thick plates.
As is clear from Table 4, all of the inventive examples B1 to B6 have macrostructures disappeared (evaluation ○ or ◎), but the comparative examples B7 to 12 all have a macro pattern, especially the comparative example. Sample symbols B11 and B12 in FIG.
[0060]
[Table 4]

[0061]
(Example 3)
This embodiment is claimed in the present invention. 2 It is related.
A thick rectangular cross-section titanium slab having an O component of 0.03 mass%, an Fe component of 0.028 mass%, and a slab thickness of 630 mm, heated by an electron beam melting method, is heated to a β region (980 ° C.), After forming the β-phase recrystallized structure by carrying out the ingot rolling with a reduction ratio of 4.2 in the β region, immediately cooling the range of 700 ° C. to the β region processing end temperature at a cooling rate of 250 ° C./hr, β → A plurality of block-rolled slabs having a thickness of 150 mm in which a fine transformation structure based on the α transformation was formed were produced.
From these piece-rolled slabs, pieces for thick plate rolling with dimensions of 150 mm thick × 1300 mm wide × 2000 mm long were collected to produce titanium thick plates based on the production conditions shown in Table 5 as sample symbols C1 to C11. .
[0062]
The rough rolling before crossing in the titanium thick plate rolling step is reduced in the plate width direction (C direction) of the block rolling slab, and the rolling direction is converted by 90 ° (the same direction as the block rolling direction: L direction). Finish rolling was performed. At that time, the finishing rolling start temperature was set to various temperatures.
As is clear from Table 5, in the VCF annealed plate, the inventive examples C1 to 6 all have the macro structure disappeared (evaluation ○ or ◎), but the comparative examples C7 to 11 all have the macro pattern. In particular, the sample symbols C10 and C11 of the comparative example had a remarkable macro pattern.
[0063]
[Table 5]

[0064]
Example 4
This embodiment is claimed in the present invention. To 3 It is related.
A rectangular cross-section titanium slab cast by an electron beam melting method with an O component of 0.04 mass%, an Fe component of 0.03 mass% and a slab thickness of 600 mm is heated to the β region (950 ° C.), Ingot rolling with a reduction ratio of 4.0 was performed, and immediately thereafter, the temperature range from the β-region heating end temperature to 700 ° C was cooled at a cooling rate of 300 ° C / hr.
A plurality of ingots with a thickness of 150 mm produced in this way were prepared, and then heated to the α region (850 ° C.), with a reduction ratio of 2.0 in the plate width direction (C direction) of the ingot slabs. Rough hot rolling is performed, and then, as finish hot rolling, cross hot rolling in which the rolling direction of the rough hot rolling is orthogonal (L direction) is used, and the cross rolling ratio is C / L between 700 to 750 ° C. = 4/10.
[0065]
Furthermore, vacuum creep plate shape correction VCF annealing was performed under conditions where the temperature range was 500 to 700 ° C. and the retention time was 1 to 20 hr. Detailed VCF annealing conditions are shown in Table 6, and the macrostructure and crystal grain size of the obtained annealed plate on the 1 mm layer surface below the surface and the 1/3 thickness plate surface below the surface were evaluated. When using a VCF annealed plate as an electrolytic drum plate, the average crystal grain size needs to be 40 μm or less, and when adding cold rolling and final annealing to the VCF annealed plate, the average crystal grain size needs to be less than 70 μm. In Table 6, when the average crystal grain size is 70 μm or more, it is described as coarse grain.
[0066]
[Table 6]

[0067]
In Table 6, the symbols D1 to D7 and D9 have a subtle degree of macro pattern (evaluation ○), but this is not good because the average crystal grain size is 70 μm or more. Symbols D21 to D28 are not good because the macro pattern remains (evaluation Δ, ×) (when the average grain size cannot be measured due to remaining rolling structure, it is described as indefinite). Symbols D8, D10 to D12, D14 to D20 are examples of the present invention in which the macro pattern is also faint (evaluation ◯) and the crystal grain size is less than 70 μm. The invention examples marked with * in Table 6 are examples in the case of producing a drum plate through a VCF annealing + cold rolling + final annealing process.
[0068]
(Example 5)
This embodiment relates to claim 5 of the present invention.
A rectangular titanium slab cast by an electron beam melting method with an O component of 0.038 mass%, an Fe component of 0.035 mass%, and a slab thickness of 630 mm is heated to the β region (950 ° C.) and reduced. A lump rolling with a ratio of 4.2 was performed, and immediately thereafter, the temperature range from the β-region heating end temperature to 700 ° C was cooled at a cooling rate of 350 ° C / hr. A plurality of pieces of 150 mm thick rolled slabs prepared in this way were prepared, and then heated to the α region (840 ° C.), with a reduction ratio of 2.2 in the plate width direction (C direction) of the divided slabs. Rough hot rolling is performed, and subsequently, as hot rolling, a cross hot rolling in which the rolling direction of the rough hot rolling is orthogonal (L direction) is performed at a temperature range of 670 ° C. to 720 ° C. with a cross rolling ratio of C / C L = 5/10 was carried out. Further, vacuum creep plate shape correction VCF annealing was performed under the conditions that the temperature range was 610 ° C. and the holding time was 5 hours.
[0069]
These VCF annealed materials are cold-rolled at a rolling rate of 20 to 60% in the longitudinal direction of the plate, annealed at a temperature range of 580 to 750 ° C., and 1 / b below the surface of the obtained annealed plate. The macrostructure and average crystal grain size on the three-plate thick layer surface were evaluated. The results are shown in Table 7.
E6 to E8 of the present invention examples have no macro pattern and an average crystal grain size of 40 μm or less, which is a good material as a titanium thick plate for an electrolytic copper foil drum.
[0070]
On the other hand, E1-E4 of the comparative examples showed a macro pattern because the cold rolling rate after annealing was less than 25% of the lower limit of the range of the present invention. In Comparative Examples E5 and E11, since the annealing temperature was lower than the range of the present invention, a macro pattern appeared. In Comparative Examples E9 and E10, since the annealing temperature is higher than the range of the present invention, the crystal grains are coarsened. In Comparative Examples E12 to E14 in which the retention time is 35 minutes, the surface oxidation is excessive, which is not good. In Comparative Examples E15 to E18, which had a high cold rolling rate of 60%, the plate shape was poor.
[0071]
[Table 7]

[0072]
【The invention's effect】
The present invention has the above-described configuration and has a restriction on equipment in hot rolling or a difficulty in production, and a 1/3 plate thickness from a plate surface layer portion by a one-step heating rolling method not based on a so-called two-step heating rolling method. It is possible to produce a titanium thick plate for an electrodeposition drum for producing an electro-deposited copper foil, which has a uniform and dense metal structure over the layer portions and does not have a macro uneven structure that causes a macro pattern (scratch defect).
[Brief description of the drawings]
BRIEF DESCRIPTION OF DRAWINGS FIG. 1 is a view showing a microstructure of an α region one-stage heating continuous unidirectionally rolled material (observation magnification 100 times).
FIG. 2 is a view showing a microstructure of an α region one-stage heating continuous cross-rolled material (observation magnification: 100 times).
FIG. 3 is a view showing a microstructure of an α region two-stage heated cross-rolled material (observation magnification: 15 times).
FIG. 4 is a view showing a microstructure of an α region two-stage heated cross-rolled material (observation magnification: 100 times).
FIG. 5 is a graph showing the relationship between the VCF annealing temperature and holding time and the microstructure average crystal grain size of the VCF annealed plate.
FIG. 6 is a diagram showing the relationship between the final annealing temperature and holding time and the microstructure average crystal grain size of the final annealing plate.

Claims

mass%,
Fe: 0.04% or less,
O: 0.1% or less, consisting of the remaining titanium and inevitable impurities, and in the surface layer portion parallel to the plate surface extending from the outermost surface to 1/3 plate thickness, the average particle size of the α phase is 40 μm or less, The surface layer is characterized in that the α phase texture is [0001] 0 ° to ± 45 ° TD and [0001] 0 ° to ± 25 ° RD (where TD is the sheet width direction and RD is the rolling direction). Titanium for copper foil production drums with excellent texture.

In producing titanium for a copper foil production drum according to claim 1, a rectangular cross-section slab having a thickness of 300 mm or more melted and cast by an electron beam melting method is heated in the β region, and a reduction ratio of 3 or more in the β region. Immediately after forming a β-phase recrystallized structure by performing ingot rolling or ingot forging, the range from the β-region working end temperature to 700 ° C. is cooled at a cooling rate of 200 ° C./hr or higher, and then heated to 880 ° C. or lower. The hot rolling is performed in the direction perpendicular to the rolling direction of the hot rolling without reheating after the hot rolling so that the cross rolling ratio is 1/10 to 6/10. The finished hot rolling is performed in a temperature range of 650 to 750 ° C., and further annealed for 1 to 10 hours in a temperature range of 570 to 680 ° C. Production method.

3. The surface layer structure according to claim 2 , wherein after the annealing, the steel sheet is further subjected to cold rolling at a rolling rate of 25 to 50%, and thereafter annealing is performed at a temperature range of 600 to 710 ° C. for 10 to 30 minutes. A method for producing titanium for copper foil production drums.