JP3870625B2 - Manufacturing method of bi-directional electrical steel sheet - Google Patents

Description

【００４４】
【発明の実施の形態】
以下に本発明の実施の形態を詳細に説明する。なお、以下に述べる％表示は質量％をあらわす。
【００４５】
（Ａ）鋼の化学組成；
Ｃ：最終焼鈍時に、脱炭に伴う（α＋γ）相からのα相への変態を利用した集合組織の制御をおこなうため、熱間圧延に供する鋼（以下、単に「素材鋼」とも記す）のＣ含有量を０．０２％以上とする。Ｃ含有量が０．０２％未満では、最終焼鈍で脱炭する前からすべてα単相となっている場合があり、変態を活用した集合組織形成ができない。Ｃ含有量が０．２０％を超えると、脱炭に長時間を要するうえ、圧延加工が困難になるので、素材鋼のＣ含有量は０．２０％以下とする。Ｃは電磁鋼板の磁気特性を大きく劣化させるので、最終製品中では少ないほどよく、多くても０．００５％以下とするのが望ましい。
【００４６】
Ｓｉ：Ｓｉは電気抵抗を増し、鉄損の一部を構成する渦電流損失を低減させる。Ｓｉはフェライト形成元素であり、Ｓｉ含有量が増すと脱炭によるα相出現の温度を高くする効果がある。本発明の製造方法では、｛１００｝面方位の形成には（α＋γ）域での高温処理が必要である。脱炭時に高温でα単相となるのが望ましいため、素材鋼のＳｉ含有量は２．４％以上とする。その含有量が４．０％を超えると鋼が脆くなり、変形抵抗が増して圧延が困難になるうえ磁束密度も低下する。これらのことから素材鋼のＳｉ含有量は４．０％以下とする。
【００４７】
Ｍｎ：Ｍｎは鋼の電気抵抗を増し、鉄損を低減させる効果がある。また、最終焼鈍時に脱炭と同時に脱Ｍｎさせることにより、｛１００｝面方位をより一層効果的に発達させることができる。このような効果を得るために、素材鋼のＭｎ含有量は０．２０％以上とする。より安定して優れた磁気特性を得るために望ましくは０．３０％以上とするのがよい。
【００４８】
Ｍｎはオーステナイト形成元素であり、２．０％を超えて含有させると脱炭に伴う（α＋γ）→α変態時に安定化されたγが残留する。残留γは非磁性であり、最終製品の磁気特性を劣化させる。これを避けるために素材鋼のＭｎ含有量は２．０％以下、望ましくは１．５％以下とする。
【００４９】
Ｍｎは最終焼鈍時に脱炭と平行して鋼板から昇華して減少するので、最終製品でのＭｎ量は、素材鋼のＣ量に依存して、すなわち脱炭に要する時間に依存して変化する。電気抵抗および鉄損の観点からは、最終製品において０．１０％以上のＭｎを含有しているのが好ましい。
【００５０】
その他の元素：
Ａｌは一般的に、鋳片の健全性確保やＮの固定などを目的として鋼に添加される場合が多く、電気抵抗を増して磁気特性を改善する効果もある。しかしながらＡｌは窒化物を形成し、脱炭最終焼鈍時の表面で酸化物を形成して｛１００｝面方位の形成を阻害して磁気特性を損なうことがある。このため、本発明においてはＡｌ含有量は少ないほどよい。多くとも０．２０％以下とするのが望ましい。
【００５１】
不可避的に混入する不純物は加工性または磁気特性を劣化させるので少ない方が好ましいが、Ｐ、Ｓ、ＮｂおよびＣｕに関しては各々０．５％以下、Ｃｒ、Ｎｉ、Ｖ、Ｗ、ＣｏおよびＭｏに関しては各々１％以下、Ｎに関しては０．０５％以下、Ｂに関しては０．００５％以下であれば含有していても本発明の効果を損なうことはない。
【００５２】
製鋼後の熱間圧延〜冷間圧延〜中間焼鈍〜冷間圧延の工程では、鋼の化学組成の変化は無視しうる程度に小さいとみなしてよく、本発明の規定する鋼の化学組成は製鋼終了後の鋼材の化学組成と同等であるとしてよい。
【００５３】
（Ｂ）圧延と中間焼鈍；
熱間圧延：熱間圧延の素材としては、鋳塊を分塊圧延したスラブ、連続鋳造によるスラブ、あるいは連続鋳造した薄鋳片などいずれでもよい。化学組成が上記の範囲を満足する鋼材は、７５０〜１０００℃の温度範囲でα＋γの２相組織となり、通常の熱間連続圧延の後段では（α＋γ）２相域での圧延となる。成分の組み合わせによっては、より高温でも２相状態になる。
【００５４】
熱間圧延時の圧延集合組織は、高温のγ相域での圧延では形成され難いが、α相域または（α＋γ）２相域での圧延では顕著に形成される。従って、化学組成が上記の範囲を満足する鋼材では、熱間圧延温度条件は特に設定しなくても、仕上げ圧延過程で圧延による集合組織の形成が可能である。また、熱間圧延後には熱間圧延集合組織を安定化させるなどの目的で熱延板焼鈍を施しても構わない。
【００５５】
冷間圧延：冷間圧延においては、圧延機のワークロール径（Ｄ）と圧延中の鋼板の厚さ（ｔ）との比（Ｄ／ｔ）が、Ｄ／ｔ≧８０またはＤ／ｔ≧１００／ｔなる関係を満足する条件で圧延するのがよい。より望ましくは、Ｄ／ｔ≧１００またはＤ／ｔ≧１００／ｔを満す条件、さらに望ましくはＤ／ｔ≧１２０またはＤ／ｔ≧１００／ｔを満す条件とするのがよい。
【００５６】
圧延の進行に応じてｔが小さくなるが、ｔの減少に応じてワークロール径Ｄを変更しても良いし、上記条件を満たす範囲内で１種類のワークロール径で最終板厚まで圧延してもよい。圧延機のロール段数や圧延速度など上記以外の条件は特に規定するものではなく、Ｄ／ｔ以外は通常使用される公知の圧延機でよい。
【００５７】
熱延鋼板を最終板厚まで冷間圧延する場合に、大圧下圧延は磁気特性にとって好ましい集合組織の形成を阻害するとして、従来は１パスあたりの圧下率を低め（１０〜１５％）に抑制する場合が多かった。しかしながらＤ／ｔが上記のような条件を満足していれば、１パスあたりの圧下率を２５％まで増大させることができる。従って本発明の方法によれば従来よりも高い圧下率での冷間圧延が可能となり、磁気特性を阻害しないで効率のよい圧延ができる。
【００５８】
しかしながら、本発明で規定するような大きなワークロール径の圧延機を用いる場合においても、１パスあたりの圧下率は、より望ましくは２０％以下、さらに望ましくは１０％以下とするのがよい。１パスあたりの圧下率が２５％以下である限り、各パスでの圧下率はパスごとに変動しても構わない。
【００５９】
冷間圧延を中間焼鈍を挟んだ複数回の圧延としておこなう場合の、それぞれの圧延における圧下率（積算圧下率、例えば一次冷間圧延であれば、一次冷間圧延の開始前の厚さと終了時の厚さから計算される圧下率）は特に限定するものではなく、通常採用される４０〜８５％の範囲であればよい。
【００６０】
冷間圧延時にはワークロールと鋼板間の摩擦係数μを０．１以下にして圧延するのが望ましい。さらに望ましくは０．０５以下がよい。摩擦係数を低下させる方法は特に限定するものではなく、圧延油や潤滑剤を使用するなど公知の方法が適用できる。圧延速度は公知の範囲でよいが、圧延速度の増大はワークロールと鋼板との摩擦係数を低める効果もあるので、圧延速度を増大させることはむしろ好ましい。
【００６１】
中間焼鈍：冷間圧延工程では圧延途中に中間焼鈍をおこなうのがよい。中間焼鈍により冷間圧延が容易になるとともに磁気特性を向上させる効果がある。最終製品の厚さが薄い場合などでは中間焼鈍を２回以上おこなってもよい。
【００６２】
中間焼鈍の均熱温度は７５０℃以上の（α＋γ）２相域がよい。より安定した磁気特性を得るには８５０℃以上とすればなおよい。鋼が（α＋γ）２相状態であれば温度は高くてもよいが、設備や操業上の限界から１２００℃程度以下とするのが好ましい。
【００６３】
中間焼鈍後の冷却はＡ１点直上から５００℃までの冷却時間を２分以下とするのがよい。均熱温度からＡ１点直上まで、および５００℃以下での冷却速度は特に限定するものではない。
【００６４】
中間焼鈍の際の加熱速度は特に限定するものではない。工業的に効率的な生産をおこなうために、急速加熱・急速冷却が可能な連続焼鈍法のような焼鈍方法を用いる場合、前記したような冷却条件となるような通板条件をそのまま適用して昇温しても、もちろん構わないし、冷却条件と別個に加熱速度を設定しても構わない。
【００６５】
均熱時間は、均熱温度が下限の７５０℃近傍である場合には数分〜数十分が望ましいが、９００℃以上の温度域で焼鈍する場合には、１０秒以上、より望ましくは３０秒以上がよい。連続焼鈍法などの工業的製造の効率化の観点から均熱時間は５分程度以下が好ましい。
【００６６】
図４は中間焼鈍における均熱係数Ｇと均熱温度、均熱時間、加熱冷却速度の関係を模式的に示すグラフである。中間焼鈍では鋼の圧延組織が整粒の再結晶組織となるように焼鈍すればよい。従ってその条件は、高温域での加熱・冷却速度が小さいほど、均熱温度が高いほど、均熱時間は短時間でよい。すなわち、均熱温度をＴ（℃）、均熱時間をｓ（秒）、７５０℃から均熱温度までの平均の加熱速度をＶｕ（℃／秒）、均熱温度から７５０℃までの平均の冷却速度をＶｄ（℃／秒）とすれば、下記式で表される均熱係数Ｇ（℃* 秒）が４５００以上となる条件で焼鈍すれば十分な再結晶組織が得られる。従ってＧ≧４５００となる範囲で諸条件を設定するのが効率的でよい。
【００６７】
【数３】

【００６８】
中間焼鈍の雰囲気は、露点を制御した水素雰囲気や窒素やアルゴン等の不活性ガス雰囲気など、非酸化性のものなら常圧あるいは減圧下のいずれでもよい。中間焼鈍は少なくともその内の１回を上述の条件でおこなえばその効果があるが、複数回中間焼鈍する場合に全ての中間焼鈍を上述の条件でおこなえばなおよい。
【００６９】
（ｃ）最終焼鈍
最終焼鈍は鋼板の形態が長尺である場合はコイル状に巻き、切板状の場合は積層して、１．３×１０⁴ Ｐａ以下の減圧下ないしは真空中でおこなう。鋼板と鋼板との間には、脱炭促進物質、もしくは脱炭促進物質と脱Ｍｎ促進物質（以下、これらを総称して「反応促進物質」とも記す）を含む焼鈍分離材を介在させて最終焼鈍する。一般的には焼鈍分離材は鋼板同士の焼付きを防止することを目的とするが、本発明では、焼鈍分離材に脱炭もしくは脱炭と脱Ｍｎを促進する機能を持たせる。なお、最終焼鈍での再結晶過程を安定化させるなどの目的で最終焼鈍前の冷間圧延鋼板に急速加熱急速冷却からなる熱処理を施しても構わない。その場合の加熱温度は中間焼鈍と同様であり、７５０℃以上の（α＋γ）２相域とするのがよく、またその上限は１２００℃程度とするのがよい。
【００７０】
脱炭促進物質としては例えばＳｉＯ₂ 、Ｃｒ₂Ｏ₃、ＴｉＯ₂ 、ＦｅＯ、Ｖ₂
Ｏ₃ 、Ｖ₂Ｏ₅、ＶＯ等の酸化物などがある。これらの酸化物は単独で使用して
もよいし、２種以上を混合して用いてもよい。これらの酸化物を鋼板表面に接触させ、減圧下で高温にすれば、酸化物が分解して放出された酸素と鋼中の炭素が反応して一酸化炭素となる等の反応により脱炭が進行するものと考えられる。反応生成 物としてのＣＯはガスとして系外に排除される。
【００７１】
脱Ｍｎ促進物質としては、最終焼鈍中に鋼板から昇華するＭｎを吸収する作用を有し、かつ、脱炭反応や鋼板の表面エネルギー状態に悪影響を及ぼさないものを用いる。このような物質としては例えばＴｉＯ₂ 、Ｔｉ₂Ｏ₃、ＳｉＯ₂ 、Ｚ
ｒＯ₂ などがある。これら物質は単独でもよいし２種以上の混合物として用いてもよい。脱炭促進物質と脱Ｍｎ促進物質とを混合して用いてもよい。
【００７２】
適切な雰囲気中では鋼板のＭｎは表面から昇華し、鋼板表面近傍にＭｎの欠乏した層（脱Ｍｎ層）が形成される。例えば脱Ｍｎ促進物質としてＴｉＯ₂ を用いる場合、ＴｉＯ₂ は鋼板から昇華したＭｎを吸収し、結合して複合酸化物（ＴｉＭｎＯ₂ ）を形成する。これにより脱Ｍｎが促進される。上記の脱Ｍｎ促進物質のうち、ＳｉＯ₂ やＴｉＯ₂ には脱炭促進作用もあるので、これら単独でも脱炭と脱Ｍｎの双方を促進することができる。
【００７３】
さらに、必須ではないが、これらの反応促進物質に加えて、高温で安定な無機物、例えば、Ａｌ₂Ｏ₃、ＣａＯ、ＺｒＯ₂ 、ＭｇＯなどの酸化物、ＳｉＣなど
の炭化物、ＢＮなどの窒化物またはホウ化物のうちの１種または２種以上を混合して含有させても構わない。これにより、反応促進物質の活性度の調整や、取り扱いを容易にするための固体状、スラリー状あるいはペースト状などへの成形が容易になり、また、鋼板への接触性が改善されるなどの効果が得られる。
【００７４】
焼鈍分離材を鋼板間に介在させる方法は任意であり、例えば粉末や液体状（スラリー状あるいはペースト状も含む）にして鋼板に塗布したり、焼鈍分離材組成物を繊維状、さらにはそれをシート状に加工したり、それらの繊維やシートにさらに粉末などを混入させたものを用いてもよい。焼鈍分離材組成物を繊維状またはさらにシート状に加工しておけば取り扱いが容易になるうえ、繊維間に生じる空隙が一酸化炭素の除去やＭｎの昇華を促進する効果も期待できるので好適である。
【００７５】
焼鈍雰囲気は減圧雰囲気ないしは真空がよく、その圧力は１．３×１０⁴ Ｐａ以下が望ましい。雰囲気の圧力が１．３×１０⁴ Ｐａを超えると一酸化炭素など反応生成物が鋼板表面から除去されにくいために反応速度が低下する。一層望ましいのは１．３×１０³ Ｐａ以下である。雰囲気圧力は低いほどよく、すなわち真空度は高いほどよいが、工業的に実施するには自ずから限界があるため、下限は１．３×１０^-3Ｐａ程度である。
【００７６】
鋼板表面への酸化物の生成や内部酸化を抑止し磁気特性の低下を避けるには全板厚にわたって脱炭が完了するまで上記減圧雰囲気で焼鈍するのがよい。しかしながら焼鈍分離材を用いて減圧下で脱炭する主たる目的は、鋼板表面に数μｍ以上の｛１００｝＜００１＞方位の再結晶粒の層を生じさせることにあるので、該再結晶粒層が生じた後は、水素を含む湿性雰囲気で、より高い圧力、ないしは常圧で脱炭しても構わない。
【００７７】
最終焼鈍では（α＋γ）２相域に均熱保持する。この温度領域での脱炭に伴う相変態により鋼板の結晶組織はα単相に変化する。均熱温度の下限は、好ましくは工業的製造が可能な脱炭速度が実現できる８５０℃以上である。その上限は、脱炭してα単相となる限りいくら高温でもよいが、１３００℃を超える高温は工業的に実現するのが困難であるので、最終焼鈍温度の上限は１３００℃程度がよい。最も効果的に｛１００｝＜００１＞方位を形成できる温度は９００〜１２００℃である。なお、鋼板表面に｛１００｝＜００１＞方位の再結晶粒の層が生じた後は、脱炭が進行する温度であれば上記のような高温でなくてもよい。
【００７８】
均熱保持時間は３０分〜１００時間の範囲とするのが望ましい。３０分未満では脱炭や脱Ｍｎが不十分で表面の｛１００｝＜００１＞方位の再結晶粒の発達が不十分であり、また鋼板の結晶粒成長も十分ではない。保持時間が１００時間を超えると焼鈍効果が飽和するうえ、結晶粒が大きくなりすぎて磁気特性が損なわれることがある。
【００７９】
最終焼鈍を終えた鋼板には、鋼板の平坦度を改善するための焼鈍や、絶縁コーティングや張力コーティング等を施すことは何ら差し支えがない。その方法は任意であり、従来、無方向性電磁鋼板や方向性電磁鋼板にて採用されているのと同様の公知の方法でよい。例えばコーティングであれば、リン酸塩系やクロム酸塩系の溶液を塗布し焼き付ける無機質系や、上記無機質系溶液にポリアクリルタイプエマルジョン等の有機樹脂を混合したものを塗布し焼き付ける有機−無機混合系のコーティングが考えられる。これらの皮膜は絶縁性を有するとともに、焼付け後の冷却時の熱収縮により板面内に等方的な張力を付加することができる。
【００８０】
【実施例】
（実施例１）
表１に示す鋼Ｃを真空鋳造し、鋳塊を熱間鍛造して８０ｍｍ厚のスラブとし、１２００℃に加熱し、熱間圧延し、酸洗して厚さ３．０ｍｍ、幅２５０ｍｍの熱延鋼板とし、次いで冷間圧延して最終厚さ０．３５ｍｍの鋼板とした。
【００８１】
【表１】

【００８２】
冷間圧延時のロール径（Ｄ）は、３８ｍｍ、６８ｍｍ、１０５ｍｍ、１５０ｍｍおよび２００ｍｍの４種類の中から適宜選択した。冷間圧延はいずれも圧延油を用いた潤滑圧延とし、中間焼鈍をしないで最終板厚まで圧延した場合と、１〜２回の中間焼鈍を挟んだ２〜３回の冷間圧延により最終板厚まで圧延した場合とについておこなった。冷間圧延時の摩擦係数μはロール周速と鋼板の出側速度との差を基に求める先進法により調査した結果０．０５〜０．２の範囲であった。１パスあたりの圧下率は全て２５％以下とした。それぞれの圧延におけるパス回数は５〜１０パスの範囲であった。
【００８３】
図５は上記圧延途中における板厚（ｔ）とＤ／ｔの関係を示すグラフである（圧延経路図とも記す）。図５で太線はＤ／ｔ＝８０またはＤ／ｔ＝１００／ｔの関係を示し、この線上および線の上部が磁気特性が良好になる範囲である。
【００８４】
中間焼鈍は連続焼鈍シミュレータを利用した。均熱温度はいずれも（α＋γ）２相域である。焼鈍後の冷却速度はＡ１点直上から５００℃までの冷却時間が２分以内の場合とそれよりも徐冷した場合の２種類とした。冷却は液体窒素ボンベから取り出した冷たい窒素ガスを鋼板に吹き付けておこなった。表２に中間焼鈍条件を示す。
【００８５】
【表２】

【００８６】
最終厚さに冷間圧延した鋼板から最終焼鈍用に長さ１００ｍｍ、幅３０ｍｍの短冊状の試験片を、図６に示すように、その長手方向が圧延方向と平行、または幅方向と平行になるように採取した。
【００８７】
焼鈍分離材として、Ａｌ₂Ｏ₃：４８質量％、ＳｉＯ₂ ：５１質量％を組成物
とする繊維化した脱炭促進物質を４０ｇ／ｍ² と、脱Ｍｎ促進作用があるＴｉＯ₂ 粉末を２０ｇ／ｍ² とを使用し、これらを試験片の間に挟んで積層した。最終焼 鈍では、０．１３Ｐａ以下の真空中で１℃／分の速度で昇温し、１０７５℃で２ ４時間保持した。焼鈍後は炉の電源を切断した炉内で冷却した。最終焼鈍後の鋼 板を化学分析した結果、Ｃ含有量は全ての試料について０．００２５％以下であ った。最終焼鈍後の各試験片の磁気特性は単板磁化特性測定装置で測定した。
【００８８】
表３に圧延経路および中間焼鈍条件と対応させて最終焼鈍後の各試験片から得られた磁気特性を示す。
【００８９】
【表３】

【００９０】
表３で英語の大文字は図５の大文字に対応するものであり、これから圧延時のワークロール径とＤ／ｔがわかる。英語の小文字は表２の中間焼鈍条件符号に対応する。例えば表３試験番号１に記載の「Ｃ→０．３５ｍｍ」は、図５に記載したように、厚さ３ｍｍの鋼板をロール径１０５ｍｍの圧延機で中間焼鈍無しで最終厚さまで圧延したことを表す。試験番号２に記載の「Ａ→Ａ1 →ｂ→Ａ3 →ａ→０．３５ｍｍ」は、厚さ３ｍｍの鋼板をロール径２００ｍｍの圧延機で１．６ｍｍに一次圧延し、条件ｂで中間焼鈍し、同一ロール径で０．７５ｍｍに２次圧延し、条件ａで中間焼鈍した後同一ロール径で最終厚さに圧延したことを表す。
【００９１】
試験番号１０に記載の「Ｂ→Ｂ2 →ｂ→Ｅ2 →０．３５ｍｍ」は、厚さ３ｍｍの鋼板をロール径１５０ｍｍの圧延機で１．６ｍｍに一次圧延し、条件ｂで中間焼鈍し、ロール径３８ｍｍの圧延機で０．３５ｍｍに２次圧延したことを表す。
【００９２】
表３からわかるように、Ｄ／ｔ≧８０またはＤ／ｔ≧１００／ｔなる関係を満す条件で冷間圧延した試験番号１〜７では、圧延方向、幅方向共に磁束密度Ｂ₁₀が１．６０以上となり、二方向性電磁鋼板として良好な磁気特性が得られた。特に中間焼鈍条件ａを挟んで冷間圧延した場合の磁気特性が良好であった。Ｄ／ｔが上記関係から外れるＥ、Ｅ1 、Ｅ2 などの圧延経路を経て冷間圧延した試験番号８〜１０では、平均磁束密度が低くなった。
【００９３】
（実施例２）
表１に示す鋼Ｂを真空鋳造して得た鋳塊を実施例１と同様の条件で厚さ３．０ｍｍに熱間圧延し、酸洗して、厚さ：０．７５ｍｍまで一次冷間圧延し、中間焼鈍した後最終圧延して厚さ：０．３５ｍｍの冷間圧延鋼板を得た。圧延は一次圧延、最終圧延共に潤滑圧延とし、１パスあたりの最大圧下率はいずれも２５％以下とした。圧延経路は、一次圧延、最終圧延共に１０５ｍｍφロールのみを用いて圧延する場合と、３ｍｍから１．２ｍｍまでを１５０ｍｍφロールで圧延し、その後３８ｍｍφロールで０．７５ｍｍに一次圧延し、その後、３８ｍｍφロールで０．３５ｍｍに最終圧延する２種類とした。図５からわかるように後者はＤ／ｔが好ましくない場合である。中間焼鈍は連続焼鈍シミュレータを利用し、均熱条件や冷却速度を種々変化させた。冷間圧延後は図６に示すように長さ１００ｍｍ、幅３０ｍｍの短冊状の磁化測定用試験片を採取し、実施例１に記載したのと同一条件で焼鈍分離剤を介在させた最終焼鈍を施した。その後、実施例１に記載したのと同様に単板磁気特性測定装置で試験片の磁気特性を測定した。表４に圧延経路、中間焼鈍条件および磁気特性測定結果をまとめて示す。
【００９４】
【表４】

【００９５】
磁気特性の幅方向での変化を見るために、圧延方向の磁束密度について、板幅方向両端部２枚（図６における「１」と「８」）の平均値から、板幅方向中央部２枚（図６における「４」と「５」）の平均値を差し引いた値（ΔＢ₁₀）も示した。幅方向のＢ₁₀平均値は、幅方向試験片３枚（図６における「９」、「１０」および「１１」）の平均値を表す。
【００９６】
表４からわかるように、750 ℃以上の (α＋γ) ２相域に加熱した後、Ａ１点真上から500 ℃までを２分以下で冷却する中間焼鈍を施した試験番号22〜30は、圧延方向、板幅方向共に、1.70Ｔを超える高い磁束密度を有し、磁束密度の板幅方向分布を示すΔＢ₁₀は0.10Ｔ以下で幅方向での磁束密度の変動も小さく、二方向性電磁鋼板として極めて良好な磁気特性を有していた。
【００９７】
中でも中間焼鈍の均熱温度が高く冷却速度も大きい試験番号２７と２８は磁束密度レベルおよびその変動共に特に良好な特性を示した。試験番号２９と３０は、中間焼鈍条件は試験番号２７、２８とほぼ同様であったが、冷間圧延時のＤ／ｔが十分には好ましい範囲でなかったために、その磁気特性は試験番号２７、２８に比較するとやや劣った結果となった。中間焼鈍の均熱温度が低すぎた試験番号２１、あるいは、中間焼鈍後のＡ１点直上から５００℃までの冷却時間が２分を超えた試験番号３１は、二方向性電磁鋼板として良好な磁気特性を有していたが、試験番号２３〜３０に比較するとやや劣った結果となった。冷間圧延時のＤ／ｔが好ましくなく、中間焼鈍をおこなわなかった試験番号３２は磁気特性が著しく劣るうえ、二方向性が実現されなかった。
【００９８】
（実施例３）
表１に示す鋼Ｂと鋼Ｄについて実施例１と同様の条件で加熱、熱間圧延、酸洗し、厚さ：０．７５ｍｍまで一次冷間圧延した。冷間圧延のワークロール径は１０５ｍｍ、Ｄ／ｔは１００／ｔ以上であった。１パスあたりの圧下率は２５％以下とした。その後、連続焼鈍シミュレータを利用し、焼鈍後の冷却速度を種々変更した以外は、表２の焼鈍条件番号ａに記載したのと同一の条件で中間焼鈍を施した。その後、ワークロール径：１０５ｍｍ、１パスあたりの圧下率は２５％以下とする条件で二次冷間圧延して厚さ：０．３５ｍｍの冷間圧延鋼板とした。
【００９９】
この冷間圧延鋼板から図６に示した方法で圧延方向または幅方向を長手方向とする長さ１００ｍｍ、幅３０ｍｍの短冊状の磁化測定用試験片を採取した。
その後、実施例１に記載したのと同一の条件で焼鈍分離材を鋼板間に介在させて積層し、０．１３Ｐａ以下の真空中で１℃／分の速度で昇温し、１０７５℃で１６時間保持する最終焼鈍を施した。冷却は焼鈍炉加熱用電源を切断した炉内で冷却した。最終焼鈍後の試験片のＣ含有量は全ての試料について０．００２５％以下であった。最終焼鈍後、各試験片の磁気特性を実施例１に記載したのと同様に単板磁気特性測定装置で測定した。
【０１００】
図７は、最終焼鈍後の試験片の圧延方向の磁束密度Ｂ₁₀に対する中間焼鈍後のＡ１点直上から５００℃までの冷却時間の影響を示すグラフであり、図７（ａ）は鋼Ｂについて、図７（ｂ）は鋼Ｄについてのものである。
【０１０１】
図８（ａ）、（ｂ）は同様にそれぞれ鋼Ｂおよび鋼Ｄの幅方向の磁束密度Ｂ₁₀に対する中間焼鈍後のＡ１点直上から５００℃までの冷却時間の影響を示すグラフである。
【０１０２】
図７および図８からわかるように、中間焼鈍後のＡ１点直上から５００℃までの冷却時間を２分以内とした場合には、鋼Ｂ、Ｄ共に１．７Ｔ以上の高い磁束密度が得られた。しかしながら、上記冷却時間が２分を超えるように徐冷した場合には、板幅方向中央部分での磁束密度が低下し、上記冷却時間が１０分であった場合にはその低下が著しく、平均磁束密度も低かった。
【０１０３】
別途、上記中間焼鈍直前および直後の微細組織を走査電子顕微鏡（ＳＥＭ）により観察し、αフェライト粒界の炭素が濃化した第二相粒の形態を調査し、同時に第二相粒にスジ状のコントラストが観察される部分の第二相粒全体に対する面積率を画像解析装置により計算した。
【０１０４】
Ａ１点の直上の温度から５００℃までの冷却時間が２分以内の場合には、αフェライト粒界に白色のコントラストをもった団塊状の残留γあるいはマルテンサイトが観察された。冷却時間が２分を超えて長い場合には、それらの領域の大部分はスジ状のコントラストが観察され、これはパーライトになっていると判断された。すなわち、中間焼鈍での冷却においてＡ１点直上から５００℃までの冷却時間を２分以内とすれば、中間焼鈍後の微細組織には残留γあるいはマルテンサイトが現れ、αフェライト粒界の炭素が濃化した第二相粒の内３０％以上が残留γあるいはマルテンサイトになっていた。２分を超えて徐冷するとαフェライト粒界にはパーライトが多量に生成し、残留γあるいはマルテンサイトの生成比率は３０％に満たなかった。
【０１０５】
（実施例４）
表１に示す鋼ＡとＢについて、実施例１と同様の方法で熱間圧延および酸洗して厚さが３．０ｍｍの熱延鋼板を得た。鋼Ａは直径１０５ｍｍのワークロールにより、鋼Ｂは直径２００ｍｍのワークロールにより、共に厚さ０．７５ｍｍまで一次冷間圧延し、いずれも表２に記載の条件ａで中間焼鈍をおこない、二次冷間圧延して最終厚さ０．３５ｍｍの冷間圧延鋼板とした。いずれの二次冷間圧延とも、１パスあたりの最大圧下率を３水準に変更して圧延した。最大圧下率を低くしたケースでは、当然のことながらパス回数は大きくなった。最終焼鈍はいずれも焼鈍分離材の適用方法を含めて実施例１に記載の条件と同一とし、実施例１に記載したのと同様の方法で最終焼鈍後の鋼板の磁気特性を測定した。
表５に得られた磁気特性を二次冷間圧延でのパス回数に対応させて示した。
【０１０６】
【表５】

【０１０７】
表５に示されているように、１パスあたりの最大圧下率が３０％に達する強圧下をおこなった試験番号４１および４５ではＢ₁₀が１．７５Ｔ以下の低い平均磁束密度しか得られなかった。１パスあたりの最大圧下率を２０％以下に低下させたものの磁気特性が良好であった。
【０１０８】
（実施例５）
表１に示す鋼Ｄについて、実施例１と同様の方法で熱間圧延および酸洗して厚さが３．０ｍｍの熱延鋼板を得た。これを直径１０５ｍｍのワークロールにより潤滑または無潤滑にて冷間圧延し、最終厚さ０．３５ｍｍの冷延鋼板を得た。冷間圧延の圧延経路は表３に記載の試番５と同一とした。１パスあたりの最大圧下率は２５％以下とした。パス回数は一次圧延が１０〜１２回、二次圧延が１０〜１１回であった。実施例１と同様の方法で最終焼鈍し、磁気特性を測定した。冷間圧延時に潤滑を施さなかった場合のワークロールと鋼板との摩擦係数μは０．２と推定された。その場合の平均磁束密度Ｂ₁₀は１．６７Ｔであった。これに対し、ワークロールと鋼板とに圧延油を塗布して一次および二次冷間圧延した場合の平均磁束密度Ｂ₁₀は１．８３Ｔであった。
【０１０９】
（実施例６）
表１に示す鋼Ａについて、実施例１と同様の条件で厚さ３．０ｍｍに熱間圧延し、酸洗して、厚さ：０．７５ｍｍまで一次冷間圧延し、中間焼鈍した後最終圧延して厚さ：０．３５ｍｍの冷間圧延鋼板を得た。圧延は一次冷間圧延、最終冷間圧延共に１０５ｍｍφロールのみを用いた潤滑圧延とし、１パスあたりの最大圧下率はいずれも２５％以下とした。中間焼鈍は連続焼鈍シミュレータを利用し、均熱条件や冷却速度を種々変化させた。冷間圧延後は実施例２と同様に短冊状の磁化測定用試験片を採取し、実施例１に記載したのと同一条件で焼鈍分離剤を介在させた最終焼鈍を施した。その後、実施例１に記載したのと同様に単板磁気特性測定装置で試験片の磁気特性を測定した。表６に圧延経路、中間焼鈍条件および磁気特性測定結果をまとめて示す。
【０１１０】
【表６】

【０１１１】
表６からわかるようにＡ１点直上から５００℃までを２分以内で冷却した試験番号６１〜６７は二方向性電磁鋼板として良好な磁気特性を有していた。しかしながら、均熱係数Ｇが４５００に満たなかった試験番号６６および６７は磁気特性がやや悪かった。Ａ１点直上から５００℃までの冷却時間が２分を超えた試験番号６８および６９は、磁気特性がよくなかった。
【０１１２】
【発明の効果】
本発明によれば、｛１００｝面が板面に平行で、圧延方向とそれに直交する方向の二方向の磁気特性に優れた二方向性電磁鋼板を安定的かつ工業的に効率よく製造することができる。従って電気機器の小型化や高効率化に大きく寄与する。
【図面の簡単な説明】
【図１】電磁鋼板の集合組織の説明図であり、同図（ａ）は｛１１０｝面が板面に平行で、＜００１＞軸が圧延方向のみに集積した組織、同図（ｂ）は｛１００｝面が板面に平行で、＜００１＞軸が板面内に特定の方向性を持たずに存在する組織、同図（ｃ）は｛１００｝面が板面に平行で、＜００１＞軸が板面内の圧延方向と幅方向に集積した組織を示す。
【図２】図２（ａ）〜（ｄ）は、冷間圧延直後の二次冷間圧延鋼板表層部の集合組織およびこれらを最終焼鈍した後の磁気特性に対する各冷間圧延条件の影響を示すグラフである。
【図３】図３（ａ）および（ｂ）は、一次冷間圧延条件をそれぞれ強圧下・無潤滑圧延と弱圧下・潤滑圧延とした場合の一次冷間圧延鋼板の厚さ方向での集合組織の変化を示すグラフである。
【図４】 中間焼鈍のヒートパターンを模式的に示すグラフである。
【図５】冷間圧延時の圧延経路を板厚ｔとＤ／ｔとの関係で示すグラフである。
【図６】磁気特性測定用の単板磁化測定用試料の採取位置を示す図である。
【図７】図７（ａ）および（ｂ）は、鋼Ｂまたは鋼Ｄの圧延方向の磁束密度Ｂ₁₀の幅方向分布例を示すグラフである。
【図８】図８（ａ）および（ｂ）は、鋼Ｂまたは鋼Ｄの幅方向の磁束密度Ｂ₁₀の幅方向分布例を示すグラフである。[0001]
BACKGROUND OF THE INVENTION
The present invention relates to a method for producing a bidirectional magnetic steel sheet having excellent magnetic properties.
[0002]
[Prior art]
Conventionally, magnetic steel sheets having a high silicon content have been used for magnetic core materials such as electric motors, generators, and transformers. This electromagnetic steel sheet is required to have a low magnetic energy loss and a high magnetic flux density in an alternating magnetic field. In order to realize these, it is effective to increase the electrical resistance of steel and form a texture in which the <001> axes of the body-centered cubic lattice, which is the direction of easy magnetization, are accumulated in the direction of the magnetic field used.
[0003]
FIG. 1 is an explanatory view of a texture of an electromagnetic steel sheet. FIG. 1 (a) shows a structure in which the {110} plane of the body-centered cubic lattice is parallel to the steel plate surface and the <001> axis is accumulated only in the rolling direction, and in the rolling direction of the steel plate like a wound core of a transformer. Suitable for applications where magnetic flux flows. An electrical steel sheet having such a texture is referred to as a unidirectional electrical steel sheet. FIG. 1B shows a structure in which the {100} plane is parallel to the steel plate surface and the <001> axis does not have a specific directionality, and various directions within the plate surface like the iron core of a rotating machine. It is suitable for an application in which magnetic flux flows through. FIG. 1C shows a structure in which the {100} plane is parallel to the steel plate surface and the <001> axis is accumulated in the rolling direction and the width direction (hereinafter also simply referred to as {100} <001> texture).
[0004]
A steel plate having a {100} <001> texture is called a bi-directional electrical steel plate because it has excellent magnetic properties in both the rolling direction and the direction perpendicular to the rolling direction. And particularly suitable for applications in which magnetic flux flows in two directions perpendicular to each other in the width direction.
[0005]
The inventors of the present invention have advanced research on an electrical steel sheet having a texture whose {100} plane is parallel to the steel sheet surface and an efficient manufacturing method thereof, and have so far disclosed techniques as described below.
[0006]
Japanese Laid-Open Patent Publication No. 7-173542 discloses a method for producing a silicon steel sheet having excellent magnetic properties, in which the {100} plane has a texture parallel to the steel sheet surface. It is a cold rolled steel sheet containing, by mass%, C: 1% or less, Si: 0.2-6.5%, Mn: 0.05-3%, in a tight coil state or a laminated state, and decarburized. It was the method of final annealing by interposing between a steel plate as an accelerator separator, or both a decarburization accelerator and a de-Mn accelerator as an annealing separator (hereinafter, this manufacturing method is called “MRD method” [Manganese Removal Decarburization Process ]).
[0007]
In the MRD method, when austenite (γ) is transformed to ferrite (α) in the decarburization process during final annealing (γ → α transformation), the recrystallized grains having {100} faces that are stable in terms of surface energy are added to the steel plate. A steel sheet having a {100} plane texture is obtained by forming the surface layer portion, and then proceeding decarburization to selectively grow the recrystallized grains. At that time, by adding a deMn-promoting substance to the annealing separator, it is possible to promote Mn sublimation from the surface of the steel sheet, thereby promoting γ → α transformation and strengthening the development of {100} plane texture. Also shown.
[0008]
In JP-A-9-20966, regarding the magnetic steel sheet by the MRD method, the ratio of the Mn concentration on the surface of the steel sheet to the Mn concentration at the center of the plate thickness is 0.90 or less, and the maximum decrease rate of the Mn concentration in the thickness direction. An electrical steel sheet having a de-Mn layer having a value of 0.05% by mass / μm or less is disclosed. The above publication also shows that a bi-directional electrical steel sheet can be obtained by subjecting a steel sheet having a final thickness to a final thickness by cold rolling a plurality of times with intermediate annealing.
[0009]
In the international patent application WO 98/20179, as a method for producing a bidirectional magnetic steel sheet having more excellent magnetic properties, a hot-rolled steel sheet containing a predetermined amount of C, Si and Mn is subjected to two or more times including intermediate annealing. It is a method of manufacturing a bidirectional magnetic steel sheet that is subjected to cold rolling by rolling, and then annealed under reduced pressure by interposing an annealing separator between the steel sheets, with rapid heating at the time of intermediate annealing A method for improving magnetic properties has been disclosed.
[0010]
[Problems to be solved by the invention]
According to the subsequent studies by the present inventors, the method disclosed in each of the above publications does not stably provide a sufficiently accumulated {100} <001> texture, and the manufacturing chance, the position in the width direction of the steel plate, etc. It has been found that the magnetic characteristics may fluctuate due to the above.
[0011]
The object of the present invention is to solve these problems and stably obtain a texture in which {100} planes parallel to the plate surface and <001> axes are accumulated in the rolling direction and the width direction. An object of the present invention is to provide a method for producing a bi-directional electrical steel sheet that is good and has little variation.
[0012]
[Means for Solving the Problems]
The present inventors have made various studies on a method for stably producing a bi-directional electrical steel sheet based on a method for preferentially developing the {100} plane orientation by the MRD method. As a result, it was found that a good {100} <001> texture can be obtained stably by limiting the cold rolling conditions and / or intermediate annealing conditions of the steel to a specific range.
[0013]
The conditions are as follows: (a) the diameter of the work roll during cold rolling is increased, the rolling reduction per pass is limited to be low, and cold rolling is performed with good lubrication; (b) during the cold rolling. At least one intermediate annealing at a temperature in the α + γ two-phase region is performed, and at least one of the intermediate annealings is performed by rapidly cooling between the point immediately above the A1 point and the pearlite transformation nose temperature during cooling. there were.
[0014]
(A) Cold rolling conditions;
Conventionally, regarding cold rolling when producing a bi-directional electrical steel sheet, for example, in the above-mentioned WO 98/20179, there is a statement that “the cold rolling ratio before and after sandwiching the intermediate annealing may be 40 to 85%”. JP-A-9-20966 discloses that “the integrated reduction ratio is 50% or more, preferably 70% or more, and that the first reduction ratio (before intermediate annealing) is 30 to 90%. It is only described as “preferred”. In the above publication, the cold rolling rate or the integrated reduction rate is defined as the value obtained from the initial thickness of the steel plate and the thickness at the time of intermediate annealing. For example, if a hot-rolled sheet having a thickness of 3 mm is cold-rolled to 0.75 mm and then subjected to intermediate annealing, then cold-rolled to 0.35 mm and subjected to final annealing, the cold-rolling before intermediate annealing is performed. The rolling rate (or rolling reduction) in rolling (primary cold rolling) is 75%, and that in cold rolling (secondary cold rolling) after intermediate annealing is 53%.
[0015]
As described above, the rolling reduction is calculated from the thicknesses before and after the start of primary cold rolling or secondary cold rolling, and is not the rolling reduction per pass described later. As described above, the prior art makes no mention of the influence of the rolling reduction per pass on the texture or magnetic properties.
[0016]
The present inventors have made various studies in order to clarify the influence of cold rolling conditions on the texture. Thickness by hot rolling a slab having a thickness of 80 mm, a width of 300 mm, and a length of 900 mm containing C: 0.066%, Si: 2.78%, Mn: 1.25% by mass%: 3mm hot-rolled steel sheet, pickled and stripped of the surface scale, then cold-rolled to a thickness of 0.75mm, subjected to intermediate annealing at 1050 ° C for 30 seconds, and then thickened Sa: Secondary cold rolled to 0.35 mm.
[0017]
FIGS. 2A to 2C are graphs showing the distribution of the X-ray integrated intensity of the surface layer portion of the steel sheet as in the secondary cold rolling and the magnetic flux density after final annealing.
FIG. 2 (a) shows the influence of the work roll diameter during cold rolling (hereinafter also simply referred to as “roll diameter”), using rolling oil for both primary cold rolling and secondary cold rolling, The maximum rolling reduction per pass is less than 30%. The roll diameter is 38 mm or 150 mm.
When rolling without using rolling oil (non-lubricating rolling), the friction coefficient μ between the work roll and the steel sheet is about 0.2, but when rolling with rolling oil (lubrication rolling) μ Is reduced to about 0.05.
[0018]
FIG. 2 (b) shows the influence of the maximum rolling reduction per pass. In both primary cold rolling and secondary cold rolling, no rolling oil is used, and the roll diameter is 150 mm. FIG. 2 (c) shows the effect of using rolling oil. In both the primary cold rolling and the secondary cold rolling, the roll diameter is 150 mm and the maximum rolling reduction per pass is 20% or less. In each figure, the vertical axis represents the diffraction intensity of each crystal orientation as a specific intensity which is a ratio to the X-ray diffraction intensity of a random sample.
[0019]
As shown in FIGS. 2A to 2C, the roll diameter is large, the rolling reduction per pass is low, and when lubricated and rolled, the {222} accumulation degree of the surface layer portion of the steel sheet increases, and the relative {200} degree of integration has decreased. When the cold-rolled steel sheet having such a rolling texture is finally annealed, a {100} <001> texture is strongly formed, and a magnetic flux density B as shown in FIG._TenIs more than 1.8T, and a bi-directional electrical steel sheet excellent in high magnetic properties is obtained.
[0020]
According to the research results of the present inventors, in order to develop the {100} <001> texture of the final product, the ratio D / t between the roll diameter D (mm) and the sheet thickness t (mm) is specified. It should be larger than the range. Moreover, it is good to roll in the range where the rolling reduction per pass does not exceed 25%. Furthermore, lubrication rolling is preferable.
[0021]
The integrated X-ray diffraction intensity of the texture of the cold-rolled steel sheet strongly depends on the rolling conditions described above as well as the integrated intensity distribution in the thickness direction of the steel sheet as well as the integrated intensity at the steel sheet surface layer.
The reason why the rolling conditions described above affect the {100} <001> texture of the final product is considered as follows.
[0022]
FIG. 3 shows a case where the same hot-rolled steel sheet as described in FIG. 2 is subjected to primary cold rolling to a thickness of 0.75 mm, and is non-lubricated and when the reduction per pass is strongly reduced. It is a graph which shows the example of the result of having investigated the texture change of the thickness direction of the primary cold-rolled steel plate at the time of carrying out a low pressure by lubrication rolling.
[0023]
In the case of rolling under high pressure without lubrication as shown in FIG. 3A, the {110} texture is strong and the {111} and {100} textures are weak on the steel sheet surface. On the other hand, when the rolling shown in FIG. 3 (b) is carried out under light pressure, {111} and {100} textures are strong not only in the surface layer but also in the thickness direction.
[0024]
In hot rolling with a large friction coefficient, it is known that a strong shear stress acts on the steel sheet surface layer portion, so that a so-called Goss orientation {110} <001> develops on the steel sheet surface layer. Also in rolling, it is suggested that a strong shear stress is acting on the steel sheet surface layer portion when rolling under high pressure without lubrication.
[0025]
It is considered that the roll diameter and the rolling reduction per pass influence the shear deformation behavior in the surface layer portion of the steel sheet and influence the rolling texture. The MRD method uses preferential growth due to the difference in surface energy of recrystallized grains having {100} planes parallel to the steel surface in the final annealing, so the texture component at the steel sheet surface layer is extremely important. is there. In the case where the shear deformation is strong in the steel sheet surface layer and the {110} texture is strong as the rolling texture, it is considered that the {100} <001> recrystallized texture is unlikely to develop in the final annealing.
[0026]
(B) Intermediate annealing conditions;
The equilibrium phase at room temperature of the steel having the chemical composition defined by the present invention is ferrite (α) and cementite (Fe_ThreeC), and a steel sheet cold-rolled through normal hot rolling exhibits a band-like structure composed of ferrite grains and cementite expanded and crushed in the rolling direction. A pearlite in which cementite is alternately layered is formed.
[0027]
When the steel plate in this state is heated to the (α + γ) two-phase region and annealed, the cementite is decomposed to be in the (α + γ) two-phase state. When this is cooled to room temperature, the γ phase is transformed, but the product at that time varies depending on the cooling conditions. When the cooling rate is high, martensite is generated, and when the cooling rate is low, pearlite is generated. It is known that the γ phase remains as residual γ when the cooling rate is further higher than the martensite generation.
[0028]
The crystal grains are in the form of being crushed and expanded in the rolling direction as they are in cold rolling, but after intermediate annealing, a recrystallized structure consisting of equiaxed grains having a grain size of about several μm to 100 μm is formed. Depending on the conditions of the intermediate annealing, a complete recrystallized structure may not be obtained, or abnormal grain growth (so-called secondary recrystallization) may occur further from normal grain growth. As described above, the form of the product from the γ phase varies in various forms depending on the cooling conditions, but most of it is distributed at the ferrite grain boundaries and hardly distributed within the ferrite grains. As described above, when the intermediate annealing condition is changed, the phase and form of the crystal structure after annealing are variously changed.
[0029]
The inventors of the present invention conducted an intermediate annealing of the Fe—Si—Mn—C based alloy under various conditions, and conducted detailed studies on the influence of the conditions on the magnetic properties of the final product. As a result, in the course of cold rolling, it is heated to the (α + γ) two-phase region of 750 ° C. or higher, and in the cooling process, a temperature at which a pearlite transformation nose occurs immediately above the point A1 (hereinafter simply referred to as “pearlite transformation nose temperature”). In the MRD method, a texture having a high final integration degree is stably formed and more uniform in the plate width direction. I learned that excellent magnetic properties can be obtained stably.
[0030]
Here, the point A1 is γ ← → (α + Fe_ThreeC) is a temperature at which the eutectoid transformation occurs, and is naturally determined from the steel composition and cooling conditions. The pearlite transformation nose temperature is a temperature at which the pearlite transformation starts in the shortest time when the steel containing γ is rapidly cooled to various temperatures and kept at that temperature to perform isothermal transformation. Although the pearlite transformation nose temperature varies depending on the chemical composition of the steel, the pearlite transformation nose temperature is generally in the range of 500 to 600 ° C. if the chemical composition is defined by the present invention. Therefore, the above cooling may be considered to be the same as cooling within 2 minutes from immediately above A1 to 500 ° C. The cooling condition is, for example, a cooling method in which the cooling rate is 2 ° C./second or more when the temperature range is cooled at a substantially constant cooling rate.
[0031]
When rapidly cooled from the two-phase region after the intermediate annealing, a part or almost all of the former γ phase is martensitic transformed and precipitates at the ferrite grain boundaries, or remains as residual γ at the ferrite grain boundaries. However, when the cooling rate is so slow that the prior γ phase does not undergo martensitic transformation, pearlite precipitates at the ferrite grain boundaries.
[0032]
When about 70% or more of the second phase enriched with carbon precipitated at the ferrite grain boundaries during cooling is pearlite, the magnetic properties of the final product deteriorate. In other words, by increasing the generation ratio of non-pearlite particles during intermediate annealing and reducing the generation ratio of pearlite to 70% or less, the variation in the magnetic properties in the plate width direction of the final product is small, and the average magnetic flux density is high. An electromagnetic steel sheet is obtained.
[0033]
The reason why the crystal structure formed by the intermediate annealing affects the cold rolling texture in the cold rolling after the intermediate annealing is presumed to be as follows. Since pearlite is softer than martensite, it easily deforms and the concentration of processing strain is reduced. On the other hand, martensite is significantly harder than the base ferrite, and strain concentrates here, changing the deformation mode and accumulated strain during cold rolling, which leads to the development of recrystallization texture in the final annealing. It is considered to have a positive effect.
[0034]
Residual γ is softer than the base ferrite, but it is significantly easier to work harden because it contains excess Si, as can be inferred from the fact that it does not undergo martensitic transformation even when brought to room temperature. It is considered that it becomes harder than the base ferrite in the initial stage of cold rolling, and has the same effect as martensite.
[0035]
In the intermediate annealing, it is only necessary that the crystal structure of the steel is a recrystallized structure of grain size in the (α + γ) two phase state. Therefore, the higher the soaking temperature is, and the shorter the heating / cooling rate in the high temperature region is, the shorter the soaking time is. That is, it is desirable that the soaking condition is intermediate annealing so that the soaking coefficient G determined from these conditions is equal to or greater than a certain limit value.
[0036]
The present invention has been completed on the basis of these newly obtained findings, and the gist thereof lies in the method for producing a bidirectional electrical steel sheet described in the following (1) to (6).
[0037]
(1) Hot rolling of steel containing C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% in mass%, Next, a method for producing a bidirectional electrical steel sheet including a step of performing cold rolling and annealing in a two-phase region under reduced pressure by interposing an annealing separator between the steel sheets, the cold rolling described above Is performed under the condition that the ratio (D / t) between the diameter D of the work roll and the steel sheet thickness t during cold rolling satisfies the relationship of D / t ≧ 80 or D / t ≧ 100 / t. A method for producing a bi-directional electrical steel sheet characterized by the above.
[0038]
(2) A steel containing, in mass%, C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% is hot-rolled and cooled. It is a method of manufacturing a bidirectional electrical steel sheet including a step of annealing in a two-phase region under reduced pressure with an annealing separator interposed between steel sheets, and the cold rolling is in the middle of the process, A bi-directional electrical steel sheet that is heated in a (α + γ) two-phase region at 750 ° C. or higher and subjected to intermediate annealing in which the cooling time from immediately above the A1 point to 500 ° C. during cooling is 2 minutes or less. Manufacturing method.
[0039]
(3) Hot rolling a steel containing C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% by mass%, A method for producing a bidirectional electrical steel sheet comprising a step of performing an intermediate rolling and annealing in a two-phase region under reduced pressure by interposing an annealing separator between steel sheets, wherein the cold rolling is performed on a work roll The ratio (D / t) between the diameter D and the steel sheet thickness t during cold rolling satisfies the relationship of D / t ≧ 80 or D / t ≧ 100 / t, and in the middle of it, 750 ° C. or higher A method for producing a bidirectional electrical steel sheet, characterized by subjecting to an (α + γ) two-phase region and performing an intermediate annealing in which the cooling time from immediately above the point A1 to 500 ° C. during cooling is 2 minutes or less .
[0040]
(4) The method for producing a bidirectional electrical steel sheet according to any one of (1) to (3) above, wherein a rolling reduction per pass of cold rolling is 25% or less.
[0041]
(5) Lubricating and rolling so that the coefficient of friction μ between the work roll and the steel sheet during cold rolling is 0.10 or less.
The manufacturing method of the bidirectional electrical steel sheet in any one of said (1)-(4).
[0042]
(6) Soaking temperature of intermediate annealing (T, ° C), soaking time (s, sec), average heating rate from 750 ° C to soaking temperature (Vu, ° C / sec) and soaking temperature to 750 ° C These conditions are selected and intermediate annealing is performed so that the soaking coefficient G calculated by the following formula is 4500 (° C. * sec) or more from the average cooling rate (Vd, ° C./sec) up to The manufacturing method of the bidirectional magnetic steel sheet in any one of said (2)-(5) to do.
[0043]
[Expression 2]

[0044]
DETAILED DESCRIPTION OF THE INVENTION
Hereinafter, embodiments of the present invention will be described in detail. In addition, the% display described below represents mass%.
[0045]
(A) Chemical composition of steel;
C: Steel used for hot rolling (hereinafter also simply referred to as “material steel”) in order to control the texture using the transformation from the (α + γ) phase to the α phase accompanying decarburization during final annealing. The C content is 0.02% or more. If the C content is less than 0.02%, it may be all α single phase before decarburization by final annealing, and texture formation utilizing transformation cannot be performed. If the C content exceeds 0.20%, it takes a long time for decarburization and rolling becomes difficult, so the C content of the material steel is set to 0.20% or less. Since C significantly deteriorates the magnetic properties of the electrical steel sheet, the lower the content in the final product, the better. It is desirable that the content be at most 0.005%.
[0046]
Si: Si increases electrical resistance and reduces eddy current loss that constitutes part of iron loss. Si is a ferrite forming element, and when the Si content increases, there is an effect of increasing the temperature at which the α phase appears due to decarburization. In the production method of the present invention, high-temperature treatment in the (α + γ) region is necessary for the formation of the {100} plane orientation. Since it is desirable to have an α single phase at high temperature during decarburization, the Si content of the raw steel is 2.4% or more. If the content exceeds 4.0%, the steel becomes brittle, the deformation resistance increases, rolling becomes difficult, and the magnetic flux density also decreases. For these reasons, the Si content of the steel material is 4.0% or less.
[0047]
Mn: Mn has the effect of increasing the electrical resistance of steel and reducing iron loss. In addition, the {100} plane orientation can be further effectively developed by de-Mn at the same time as decarburization during the final annealing. In order to obtain such an effect, the Mn content of the material steel is set to 0.20% or more. In order to obtain more stable and excellent magnetic characteristics, it is desirable that the content be 0.30% or more.
[0048]
Mn is an austenite-forming element, and if it is contained in an amount exceeding 2.0%, γ stabilized during (α + γ) → α transformation associated with decarburization remains. Residual γ is non-magnetic and degrades the magnetic properties of the final product. In order to avoid this, the Mn content of the raw steel is 2.0% or less, preferably 1.5% or less.
[0049]
Since Mn sublimates from the steel sheet and decreases in parallel with decarburization at the time of final annealing, the amount of Mn in the final product changes depending on the C amount of the raw steel, that is, depending on the time required for decarburization. . From the viewpoint of electrical resistance and iron loss, the final product preferably contains 0.10% or more of Mn.
[0050]
Other elements:
In general, Al is often added to steel for the purpose of securing soundness of slabs or fixing N, and has an effect of increasing the electric resistance and improving the magnetic properties. However, Al forms a nitride, and an oxide is formed on the surface at the time of final decarburization annealing, thereby obstructing the formation of {100} plane orientation and impairing magnetic properties. For this reason, in the present invention, the lower the Al content, the better. At most 0.20% or less is desirable.
[0051]
Impurities that are inevitably mixed are preferably less because they degrade workability or magnetic properties. However, for P, S, Nb, and Cu, each is 0.5% or less, and for Cr, Ni, V, W, Co, and Mo, respectively. The content of each is 1% or less, N is 0.05% or less, and B is 0.005% or less. Even if it is contained, the effect of the present invention is not impaired.
[0052]
In the processes of hot rolling, cold rolling, intermediate annealing, and cold rolling after steelmaking, changes in the chemical composition of the steel may be considered to be negligibly small, and the chemical composition of the steel defined in the present invention is steelmaking. It may be equivalent to the chemical composition of the steel material after completion.
[0053]
(B) rolling and intermediate annealing;
Hot rolling: The raw material for hot rolling may be a slab obtained by rolling an ingot, a slab obtained by continuous casting, or a thin cast piece obtained by continuous casting. A steel material having a chemical composition satisfying the above range has an α + γ two-phase structure in the temperature range of 750 to 1000 ° C., and is rolled in the (α + γ) two-phase region after the normal hot continuous rolling. Depending on the combination of the components, a two-phase state is obtained even at a higher temperature.
[0054]
The rolling texture at the time of hot rolling is hardly formed by rolling in a high-temperature γ phase region, but is remarkably formed by rolling in an α phase region or (α + γ) 2 phase region. Therefore, in a steel material having a chemical composition satisfying the above range, a texture can be formed by rolling in the finish rolling process even if the hot rolling temperature condition is not particularly set. In addition, after hot rolling, hot rolled sheet annealing may be performed for the purpose of stabilizing the hot rolling texture.
[0055]
Cold rolling: In cold rolling, the ratio (D / t) between the work roll diameter (D) of the rolling mill and the thickness (t) of the steel sheet during rolling is D / t ≧ 80 or D / t ≧ It is good to roll on the conditions which satisfy | fill the relationship of 100 / t. More preferably, the conditions satisfying D / t ≧ 100 or D / t ≧ 100 / t, and more preferably the conditions satisfying D / t ≧ 120 or D / t ≧ 100 / t.
[0056]
Although t decreases as the rolling progresses, the work roll diameter D may be changed as t decreases, and the roll is rolled to the final thickness with one work roll diameter within a range that satisfies the above conditions. May be. Conditions other than those described above, such as the number of roll stages and rolling speed of the rolling mill, are not particularly specified, and other known rolling mills that are usually used may be used except for D / t.
[0057]
When cold-rolling hot-rolled steel sheets to the final thickness, large reduction rolling has conventionally suppressed the reduction rate per pass to be low (10 to 15%), as it inhibits formation of a texture favorable for magnetic properties. There were many cases to do. However, if D / t satisfies the above conditions, the rolling reduction per pass can be increased to 25%. Therefore, according to the method of the present invention, it is possible to perform cold rolling at a higher rolling reduction than before, and to perform efficient rolling without impairing magnetic properties.
[0058]
However, even when a rolling mill having a large work roll diameter as defined in the present invention is used, the rolling reduction per pass is more desirably 20% or less, and further desirably 10% or less. As long as the rolling reduction per pass is 25% or less, the rolling reduction in each pass may vary from pass to pass.
[0059]
When performing cold rolling as multiple rollings with intermediate annealing, the rolling reduction in each rolling (integrated rolling reduction, for example, if it is primary cold rolling, the thickness before the start of primary cold rolling and at the end The rolling reduction calculated from the thickness of the film is not particularly limited, and may be in the range of 40 to 85% that is usually employed.
[0060]
In cold rolling, it is desirable to perform rolling with a coefficient of friction μ between the work roll and the steel sheet being 0.1 or less. Furthermore, 0.05 or less is desirable. The method for reducing the friction coefficient is not particularly limited, and a known method such as using rolling oil or a lubricant can be applied. Although the rolling speed may be in a known range, increasing the rolling speed also has the effect of lowering the friction coefficient between the work roll and the steel sheet, so it is preferable to increase the rolling speed.
[0061]
Intermediate annealing: In the cold rolling process, intermediate annealing is preferably performed during rolling. Intermediate annealing has the effect of facilitating cold rolling and improving magnetic properties. When the final product is thin, intermediate annealing may be performed twice or more.
[0062]
The soaking temperature of the intermediate annealing is preferably an (α + γ) two-phase region of 750 ° C. or higher. In order to obtain more stable magnetic characteristics, the temperature should be 850 ° C. or higher. If the steel is in the (α + γ) two-phase state, the temperature may be high, but is preferably about 1200 ° C. or less because of equipment and operational limits.
[0063]
For cooling after the intermediate annealing, the cooling time from immediately above the A1 point to 500 ° C. should be 2 minutes or less. The cooling rate from the soaking temperature to just above the A1 point and at 500 ° C. or lower is not particularly limited.
[0064]
The heating rate during the intermediate annealing is not particularly limited. In order to perform industrially efficient production, when using an annealing method such as a continuous annealing method capable of rapid heating / cooling, apply the sheet passing conditions as described above as they are. Of course, the temperature may be raised or the heating rate may be set separately from the cooling condition.
[0065]
The soaking time is preferably several minutes to several tens of minutes when the soaking temperature is near the lower limit of 750 ° C., but is 10 seconds or more, more preferably 30 when annealing is performed in a temperature range of 900 ° C. or more. More than a second is good. From the viewpoint of increasing the efficiency of industrial production such as continuous annealing, the soaking time is preferably about 5 minutes or less.
[0066]
FIG. 4 is a graph schematically showing the relationship between the soaking coefficient G, soaking temperature, soaking time, and heating / cooling rate in the intermediate annealing. In the intermediate annealing, annealing may be performed so that the rolled structure of the steel becomes a recrystallized structure of sized particles. Therefore, the conditions are such that the shorter the heating / cooling rate in the high temperature range, the higher the soaking temperature, and the shorter the soaking time. That is, the soaking temperature is T (° C.), the soaking time is s (seconds), the average heating rate from 750 ° C. to the soaking temperature is Vu (° C./sec), and the average soaking temperature is 750 ° C. If the cooling rate is Vd (° C./second), a sufficient recrystallized structure can be obtained if annealing is performed under the condition that the soaking coefficient G (° C. * second) represented by the following formula is 4500 or more. Therefore, it is efficient to set various conditions in a range where G ≧ 4500.
[0067]
[Equation 3]

[0068]
The atmosphere of the intermediate annealing may be normal pressure or reduced pressure as long as it is non-oxidizing, such as a hydrogen atmosphere with a controlled dew point or an inert gas atmosphere such as nitrogen or argon. The intermediate annealing is effective if at least one of them is performed under the above-described conditions, but it is more preferable if all the intermediate annealing is performed under the above-mentioned conditions when performing intermediate annealing a plurality of times.
[0069]
(C) Final annealing
In the final annealing, when the form of the steel sheet is long, it is wound in a coil shape, and in the case of a cut sheet shape, it is laminated, and 1.3 × 10^FourThe reaction is performed under a reduced pressure of Pa or less or in a vacuum. The steel sheet is interleaved with an annealing separator containing a decarburization promoting substance, or a decarburization promoting substance and a deMn promoting substance (hereinafter collectively referred to as “reaction promoting substance”). Annealing. In general, the annealing separator aims to prevent seizure between steel plates, but in the present invention, the annealing separator has a function of promoting decarburization or decarburization and de-Mn. In addition, you may perform the heat processing which consists of rapid heating rapid cooling to the cold rolled steel plate before final annealing for the purpose of stabilizing the recrystallization process in final annealing. In this case, the heating temperature is the same as that in the intermediate annealing, and the (α + γ) two-phase region of 750 ° C. or higher is preferable, and the upper limit is preferably about 1200 ° C.
[0070]
As a decarburization promoting substance, for example, SiO₂, Cr₂O_ThreeTiO₂, FeO, V₂
O_Three, V₂O_FiveAnd oxides such as VO. Use these oxides alone
Alternatively, two or more kinds may be mixed and used. If these oxides are brought into contact with the surface of the steel sheet and heated to a high temperature under reduced pressure, decarburization is caused by a reaction such as the decomposition of the oxide and the release of oxygen and the carbon in the steel to react with carbon monoxide. It is considered to progress. CO as a reaction product is excluded from the system as a gas.
[0071]
As the de-Mn promoting substance, a substance that has an action of absorbing Mn sublimated from the steel sheet during the final annealing and does not adversely affect the decarburization reaction or the surface energy state of the steel sheet is used. As such a substance, for example, TiO₂, Ti₂O_Three, SiO₂, Z
rO₂and so on. These substances may be used alone or as a mixture of two or more. A mixture of a decarburization promoting substance and a deMn promoting substance may be used.
[0072]
In an appropriate atmosphere, Mn of the steel sheet sublimates from the surface, and a Mn-deficient layer (de-Mn layer) is formed in the vicinity of the steel sheet surface. For example, TiO as a deMn promoting substance₂When using TiO₂Absorbs Mn sublimated from the steel sheet and bonds it to form a composite oxide (TiMnO₂). This promotes removal of Mn. Among the above de-Mn promoting substances, SiO₂And TiO₂Since there is also a decarburization promoting action, both these can promote both decarburization and deMn.
[0073]
Furthermore, although not essential, in addition to these reaction promoters, inorganic substances that are stable at high temperatures, such as Al₂O_Three, CaO, ZrO₂, Oxides such as MgO, SiC, etc.
One kind or two or more kinds of nitrides or borides such as carbides, BN and the like may be mixed and contained. This makes it easy to adjust the activity of the reaction promoting substance, to form a solid, slurry or paste for easy handling, and to improve the contact with the steel sheet. An effect is obtained.
[0074]
The method of interposing the annealing separator between the steel plates is arbitrary. For example, it is applied to the steel plate in the form of powder or liquid (including slurry or paste), the annealing separator composition is fibrous, and further It may be processed into a sheet shape, or those obtained by further mixing powder or the like into those fibers or sheets. If the annealing separator composition is processed into a fiber or a sheet, it is easy to handle, and the void generated between the fibers can be expected to promote the removal of carbon monoxide and the sublimation of Mn. is there.
[0075]
The annealing atmosphere is preferably a reduced pressure atmosphere or a vacuum, and the pressure is 1.3 × 10^FourPa or less is desirable. Atmospheric pressure is 1.3 × 10^FourIf it exceeds Pa, the reaction rate is lowered because reaction products such as carbon monoxide are difficult to be removed from the steel sheet surface. More desirable 1.3 × 10^ThreePa or less. The lower the atmospheric pressure, the better, the better the degree of vacuum, but the lower limit is 1.3 × 10 because there is a natural limit to industrial implementation.^-3It is about Pa.
[0076]
In order to suppress the formation of oxides on the surface of the steel sheet and internal oxidation and to avoid deterioration of the magnetic properties, it is preferable to anneal in the reduced pressure atmosphere until decarburization is completed over the entire thickness. However, the main purpose of decarburizing using an annealing separator under reduced pressure is to generate a layer of recrystallized grains with a {100} <001> orientation of several μm or more on the steel sheet surface. After this occurs, decarburization may be performed at a higher pressure or normal pressure in a wet atmosphere containing hydrogen.
[0077]
In the final annealing, soaking is maintained in the (α + γ) two-phase region. Due to the phase transformation accompanying decarburization in this temperature range, the crystal structure of the steel sheet changes to α single phase. The lower limit of the soaking temperature is preferably 850 ° C. or higher at which a decarburization rate capable of industrial production can be realized. The upper limit may be as high as possible as long as it decarburizes and becomes an α single phase, but since it is difficult to industrially realize a high temperature exceeding 1300 ° C., the upper limit of the final annealing temperature is preferably about 1300 ° C. The temperature at which {100} <001> orientation can be most effectively formed is 900 to 1200 ° C. In addition, after the layer of the recrystallized grains having the {100} <001> orientation is generated on the steel plate surface, the above-described high temperature is not necessary as long as the decarburization proceeds.
[0078]
The soaking time is preferably in the range of 30 minutes to 100 hours. If it is less than 30 minutes, decarburization and deMn removal are insufficient, and the development of recrystallized grains in the {100} <001> orientation on the surface is insufficient, and the crystal grain growth of the steel sheet is not sufficient. If the holding time exceeds 100 hours, the annealing effect is saturated and the crystal grains become too large, and the magnetic properties may be impaired.
[0079]
The steel sheet that has been subjected to the final annealing may be subjected to any annealing for improving the flatness of the steel sheet, or an insulating coating or a tension coating. The method is arbitrary, and may be a known method similar to that conventionally employed for non-oriented electrical steel sheets and directional electrical steel sheets. For example, in the case of a coating, an inorganic system in which a phosphate-based or chromate-based solution is applied and baked, or an organic-inorganic mixture in which an organic resin such as a polyacryl type emulsion is mixed and baked in the above inorganic solution. System coatings are conceivable. These coatings have insulating properties and can apply isotropic tension in the plate surface due to thermal contraction during cooling after baking.
[0080]
【Example】
Example 1
Steel C shown in Table 1 is vacuum cast, and the ingot is hot forged into a slab of 80 mm thickness, heated to 1200 ° C., hot rolled, pickled and heated to a thickness of 3.0 mm and a width of 250 mm. The steel sheet was then rolled and then cold-rolled to a steel sheet having a final thickness of 0.35 mm.
[0081]
[Table 1]

[0082]
The roll diameter (D) at the time of cold rolling was appropriately selected from 4 types of 38 mm, 68 mm, 105 mm, 150 mm and 200 mm. Cold rolling is lubricated rolling using rolling oil, and the final plate is obtained by rolling to the final thickness without intermediate annealing and by 2-3 cold rollings with 1 to 2 intermediate annealings. This was done for the case of rolling to thickness. The coefficient of friction μ during cold rolling was in the range of 0.05 to 0.2 as a result of investigation by an advanced method obtained on the basis of the difference between the roll peripheral speed and the exit speed of the steel sheet. The rolling reduction per pass was 25% or less. The number of passes in each rolling ranged from 5 to 10 passes.
[0083]
FIG. 5 is a graph showing the relationship between sheet thickness (t) and D / t during the rolling (also referred to as a rolling path diagram). In FIG. 5, the thick line indicates a relationship of D / t = 80 or D / t = 100 / t, and the area on the line and the upper part of the line is a range in which the magnetic characteristics are good.
[0084]
For the intermediate annealing, a continuous annealing simulator was used. The soaking temperature is in the (α + γ) two-phase region. The cooling rate after annealing was set to two types when the cooling time from directly above the A1 point to 500 ° C. was within 2 minutes and when the cooling time was gradually lowered. Cooling was performed by blowing cold nitrogen gas taken out from the liquid nitrogen cylinder onto the steel plate. Table 2 shows the intermediate annealing conditions.
[0085]
[Table 2]

[0086]
A strip-shaped test piece having a length of 100 mm and a width of 30 mm for a final annealing from a steel sheet cold-rolled to a final thickness, as shown in FIG. 6, its longitudinal direction is parallel to the rolling direction or parallel to the width direction. It collected so that it might become.
[0087]
As an annealing separator, Al₂O_Three: 48% by mass, SiO₂: 51% by mass of composition
40 g / m of the decarburization promoting substance made into fiber²And TiO having a de-Mn promoting action₂  20 g / m of powder²And these were laminated between test pieces. In the final annealing, the temperature was raised at a rate of 1 ° C./min in a vacuum of 0.13 Pa or less and held at 1075 ° C. for 24 hours. After annealing, it was cooled in a furnace where the power source of the furnace was cut off. As a result of chemical analysis of the steel plate after the final annealing, the C content was 0.0025% or less for all samples. The magnetic properties of each test piece after final annealing were measured with a single plate magnetization property measuring device.
[0088]
Table 3 shows the magnetic properties obtained from the test pieces after the final annealing in correspondence with the rolling path and the intermediate annealing conditions.
[0089]
[Table 3]

[0090]
In Table 3, uppercase letters in English correspond to uppercase letters in FIG. 5, and from this, the work roll diameter and D / t during rolling can be found. The English lower case letters correspond to the intermediate annealing condition codes in Table 2. For example, “C → 0.35 mm” described in Test No. 1 in Table 3 indicates that, as shown in FIG. 5, a steel plate having a thickness of 3 mm was rolled to the final thickness without intermediate annealing with a rolling mill having a roll diameter of 105 mm. To express. “A → A1 → b → A3 → a → 0.35 mm” described in Test No. 2 is a primary rolling of a steel plate with a thickness of 3 mm to 1.6 mm with a rolling mill with a roll diameter of 200 mm and intermediate annealing under condition b. , Secondary rolling to 0.75 mm with the same roll diameter, intermediate annealing under condition a, and then rolling to the final thickness with the same roll diameter.
[0091]
“B → B 2 → b → E 2 → 0.35 mm” described in Test No. 10 is a method in which a steel sheet having a thickness of 3 mm is primarily rolled to 1.6 mm by a rolling mill having a roll diameter of 150 mm, subjected to intermediate annealing under condition b, It represents that the secondary rolling was performed to 0.35 mm with a rolling mill having a diameter of 38 mm.
[0092]
As can be seen from Table 3, in Test Nos. 1 to 7 that were cold-rolled under the condition that D / t ≧ 80 or D / t ≧ 100 / t, the magnetic flux density B in both the rolling direction and the width direction._TenWas 1.60 or more, and good magnetic properties were obtained as a bi-directional electrical steel sheet. In particular, the magnetic properties were good when cold-rolled with the intermediate annealing condition a interposed therebetween. In test numbers 8 to 10 that were cold-rolled through rolling paths such as E, E1, and E2 where D / t deviated from the above relationship, the average magnetic flux density was low.
[0093]
(Example 2)
An ingot obtained by vacuum casting steel B shown in Table 1 was hot-rolled to a thickness of 3.0 mm under the same conditions as in Example 1, pickled, and first cold to a thickness of 0.75 mm. After rolling and intermediate annealing, final rolling was performed to obtain a cold rolled steel sheet having a thickness of 0.35 mm. Rolling was lubricated rolling for both primary rolling and final rolling, and the maximum rolling reduction per pass was 25% or less. The rolling path is the case of rolling using only 105 mmφ rolls for both primary rolling and final rolling, and rolling from 3 mm to 1.2 mm with 150 mmφ rolls, then primary rolling to 0.75 mm with 38 mmφ rolls, and then 38 mmφ rolls Thus, two types of final rolling to 0.35 mm were made. As can be seen from FIG. 5, the latter is a case where D / t is not preferable. For intermediate annealing, a continuous annealing simulator was used, and soaking conditions and cooling rates were varied. After cold rolling, as shown in FIG. 6, a strip-shaped test piece for magnetization measurement having a length of 100 mm and a width of 30 mm was collected and subjected to final annealing in which an annealing separator was interposed under the same conditions as described in Example 1. Was given. Thereafter, in the same manner as described in Example 1, the magnetic properties of the test piece were measured with a single plate magnetic property measuring apparatus. Table 4 summarizes the rolling path, intermediate annealing conditions, and magnetic property measurement results.
[0094]
[Table 4]

[0095]
In order to see the change in the magnetic property in the width direction, with respect to the magnetic flux density in the rolling direction, from the average value of the two end portions in the plate width direction (“1” and “8” in FIG. 6), the center portion 2 in the plate width direction. The value (ΔB) obtained by subtracting the average value of the sheets (“4” and “5” in FIG. 6)_Ten). B in the width direction_TenThe average value represents the average value of three test pieces in the width direction (“9”, “10”, and “11” in FIG. 6).
[0096]
As can be seen from Table 4, the test numbers 22-30, which were heated to a (α + γ) two-phase region of 750 ° C or higher and then subjected to intermediate annealing for cooling from just above A1 to 500 ° C in 2 minutes or less, were rolled ΔB, which has a high magnetic flux density exceeding 1.70 T in both the direction and the plate width direction, and indicates the plate width direction distribution of the magnetic flux density_TenIs less than 0.10T and small fluctuation of magnetic flux density in the width directionBloomAs a bi-directional electrical steel sheet, it had extremely good magnetic properties.
[0097]
In particular, Test Nos. 27 and 28, which have a high soaking temperature in the intermediate annealing and a high cooling rate, showed particularly good characteristics both in terms of the magnetic flux density level and its fluctuation. Test Nos. 29 and 30 were substantially the same as the test Nos. 27 and 28 in the intermediate annealing conditions, but the D / t during cold rolling was not a preferable range. Compared with 28, the result was slightly inferior. Test No. 21 in which the soaking temperature of the intermediate annealing is too low, or Test No. 31 in which the cooling time from immediately above the A1 point to 500 ° C. after the intermediate annealing exceeds 2 minutes is a good magnetic property as a bidirectional magnetic steel sheet. Although it had the characteristics, the results were slightly inferior to those of test numbers 23 to 30. D / t at the time of cold rolling was not preferable, and test number 32 in which the intermediate annealing was not performed was remarkably inferior in magnetic properties, and bi-directionality was not realized.
[0098]
(Example 3)
Steel B and steel D shown in Table 1 were heated, hot-rolled and pickled under the same conditions as in Example 1 and subjected to primary cold rolling to a thickness of 0.75 mm. The cold-rolled work roll diameter was 105 mm and D / t was 100 / t or more. The rolling reduction per pass was 25% or less. Then, the intermediate annealing was performed on the same conditions as described in the annealing condition number a of Table 2 except having changed variously the cooling rate after annealing using a continuous annealing simulator. Thereafter, a cold rolled steel sheet having a thickness of 0.35 mm was obtained by secondary cold rolling under a condition that the work roll diameter was 105 mm and the rolling reduction per pass was 25% or less.
[0099]
From this cold rolled steel sheet, a strip-shaped test piece for magnetization measurement having a length of 100 mm and a width of 30 mm with the rolling direction or the width direction as the longitudinal direction was collected by the method shown in FIG.
Thereafter, an annealing separator was laminated between the steel plates under the same conditions as described in Example 1, and the temperature was increased at a rate of 1 ° C./min in a vacuum of 0.13 Pa or less, and 16 ° C. at 1675 ° C. The final annealing which hold | maintains time was given. Cooling was performed in a furnace in which the power source for heating the annealing furnace was cut off. The C content of the test piece after the final annealing was 0.0025% or less for all the samples. After the final annealing, the magnetic properties of each test piece were measured with a single plate magnetic property measuring apparatus in the same manner as described in Example 1.
[0100]
FIG. 7 shows the magnetic flux density B in the rolling direction of the test piece after the final annealing._Ten7A and 7B are graphs showing the influence of the cooling time from immediately above A1 point to 500 ° C. after intermediate annealing, FIG. 7A is for steel B, and FIG.
[0101]
FIGS. 8A and 8B show the magnetic flux density B in the width direction of steel B and steel D, respectively._TenIt is a graph which shows the influence of the cooling time from A1 point immediately after intermediate annealing to 500 degreeC with respect to.
[0102]
As can be seen from FIGS. 7 and 8, when the cooling time from immediately above the A1 point after intermediate annealing to 500 ° C. is within 2 minutes, a high magnetic flux density of 1.7 T or more is obtained for both steel B and D. It was. However, when the cooling time is gradually cooled so as to exceed 2 minutes, the magnetic flux density in the central portion in the plate width direction is reduced, and when the cooling time is 10 minutes, the reduction is significant, and the average The magnetic flux density was also low.
[0103]
Separately, the microstructure immediately before and immediately after the intermediate annealing was observed with a scanning electron microscope (SEM), and the morphology of the second phase grains in which the carbon of the α ferrite grain boundaries was concentrated was investigated. The area ratio with respect to the entire second phase grains in the part where the contrast of the film was observed was calculated by an image analyzer.
[0104]
When the cooling time from the temperature immediately above the point A1 to 500 ° C. was within 2 minutes, nodular residual γ or martensite having a white contrast was observed at the α ferrite grain boundary. When the cooling time was longer than 2 minutes, a streak-like contrast was observed in most of these regions, which was judged to be pearlite. That is, if the cooling time from immediately above the A1 point to 500 ° C. within cooling in the intermediate annealing is within 2 minutes, residual γ or martensite appears in the microstructure after the intermediate annealing, and carbon in the α ferrite grain boundary is concentrated. More than 30% of the converted second phase grains were residual γ or martensite. When it was slowly cooled for more than 2 minutes, a large amount of pearlite was produced at the α ferrite grain boundary, and the production rate of residual γ or martensite was less than 30%.
[0105]
(Example 4)
Steels A and B shown in Table 1 were hot-rolled and pickled in the same manner as in Example 1 to obtain hot-rolled steel sheets having a thickness of 3.0 mm. Steel A was first cold-rolled to a thickness of 0.75 mm by a work roll having a diameter of 105 mm, and steel B was a work roll having a diameter of 200 mm, both of which were subjected to intermediate annealing under the condition a described in Table 2, Cold rolling was performed to obtain a cold rolled steel sheet having a final thickness of 0.35 mm. In any secondary cold rolling, the maximum rolling reduction per pass was changed to 3 levels and rolled. As a matter of course, the number of passes increased in the case where the maximum rolling reduction was lowered. The final annealing was the same as that described in Example 1 including the application method of the annealing separator, and the magnetic properties of the steel sheet after final annealing were measured in the same manner as described in Example 1.
Table 5 shows the magnetic characteristics obtained in correspondence with the number of passes in secondary cold rolling.
[0106]
[Table 5]

[0107]
As shown in Table 5, in test Nos. 41 and 45 in which the maximum reduction rate per pass reaches 30%,_TenHowever, only a low average magnetic flux density of 1.75 T or less was obtained. Although the maximum rolling reduction per pass was reduced to 20% or less, the magnetic properties were good.
[0108]
(Example 5)
Steel D shown in Table 1 was hot-rolled and pickled in the same manner as in Example 1 to obtain a hot-rolled steel sheet having a thickness of 3.0 mm. This was cold-rolled with a work roll having a diameter of 105 mm with or without lubrication to obtain a cold-rolled steel sheet with a final thickness of 0.35 mm. The rolling path of the cold rolling was the same as the trial number 5 shown in Table 3. The maximum rolling reduction per pass was 25% or less. The number of passes was 10 to 12 times for primary rolling and 10 to 11 times for secondary rolling. Final annealing was performed in the same manner as in Example 1, and the magnetic properties were measured. The coefficient of friction μ between the work roll and the steel plate when lubrication was not performed during cold rolling was estimated to be 0.2. Average magnetic flux density B in that case_TenWas 1.67T. On the other hand, the average magnetic flux density B when the rolling oil is applied to the work roll and the steel plate and primary and secondary cold rolling are performed._TenWas 1.83T.
[0109]
(Example 6)
Steel A shown in Table 1 was hot-rolled to a thickness of 3.0 mm under the same conditions as in Example 1, pickled, first cold-rolled to a thickness of 0.75 mm, and finally subjected to intermediate annealing. A cold rolled steel sheet having a thickness of 0.35 mm was obtained by rolling. Rolling was lubricated rolling using only 105 mmφ rolls for both the primary cold rolling and the final cold rolling, and the maximum rolling reduction per pass was 25% or less. For intermediate annealing, a continuous annealing simulator was used, and soaking conditions and cooling rates were varied. After cold rolling, strip-like test pieces for magnetization measurement were collected in the same manner as in Example 2 and subjected to final annealing with an annealing separator in the same conditions as described in Example 1. Thereafter, in the same manner as described in Example 1, the magnetic properties of the test piece were measured with a single plate magnetic property measuring apparatus. Table 6 summarizes the rolling path, intermediate annealing conditions, and magnetic property measurement results.
[0110]
[Table 6]

[0111]
As can be seen from Table 6, Test Nos. 61 to 67 which were cooled within 2 minutes from just above A1 to 500 ° C. had good magnetic properties as bidirectional magnetic steel sheets. However, test numbers 66 and 67 having a soaking coefficient G of less than 4500 had slightly poor magnetic properties. Test Nos. 68 and 69 in which the cooling time from directly above the A1 point to 500 ° C. exceeded 2 minutes had poor magnetic properties.
[0112]
【The invention's effect】
According to the present invention, a bi-directional electrical steel sheet having a {100} plane parallel to a plate surface and excellent in magnetic properties in two directions, ie, a rolling direction and a direction perpendicular thereto, can be stably and industrially produced efficiently. Can do. Therefore, it greatly contributes to miniaturization and high efficiency of electric equipment.
[Brief description of the drawings]
FIG. 1 is an explanatory diagram of a texture of an electromagnetic steel sheet, in which FIG. (A) is a structure in which the {110} plane is parallel to the plate surface and the <001> axis is accumulated only in the rolling direction, FIG. Is a structure in which the {100} plane is parallel to the plate surface and the <001> axis does not have a specific direction in the plate surface, and FIG. 10 (c) shows that the {100} plane is parallel to the plate surface, The <001> axis shows a structure accumulated in the rolling direction and the width direction in the plate surface.
FIGS. 2 (a) to 2 (d) show the effect of each cold rolling condition on the texture of the secondary cold rolled steel sheet surface layer immediately after cold rolling and the magnetic properties after final annealing. It is a graph to show.
FIGS. 3 (a) and 3 (b) show the aggregation in the thickness direction of the primary cold-rolled steel sheet when the primary cold rolling conditions are strong rolling / non-lubricating rolling and weak rolling / lubricating rolling, respectively. It is a graph which shows the change of an organization.
[Fig. 4] Schematic of heat pattern of intermediate annealingShown inIt is a graph.
FIG. 5 is a graph showing a rolling path at the time of cold rolling in relation to a sheet thickness t and D / t.
FIG. 6 is a diagram showing a sampling position of a single plate magnetization measurement sample for measuring magnetic properties.
7 (a) and 7 (b) show the magnetic flux density B in the rolling direction of steel B or steel D. FIG._TenIt is a graph which shows the example of distribution of the width direction.
8 (a) and 8 (b) show the magnetic flux density B in the width direction of steel B or steel D. FIG._TenIt is a graph which shows the example of distribution of the width direction.

Claims (6)

Heat of steel containing, by mass%, C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% , the balance being Fe and impurities A method for producing a bidirectional electrical steel sheet comprising a step of performing cold rolling, followed by cold rolling, and annealing in a (α + γ) two-phase region under reduced pressure by interposing an annealing separator between the steel plates. In the cold rolling, the ratio (D / t) between the diameter D (mm) of the work roll and the steel sheet thickness t (mm) during the cold rolling is D / t ≧ 80 or D / t ≧ 100 / A method for producing a bi-directional electrical steel sheet, which is performed under a condition satisfying a relationship t. The steel containing, by mass%, C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% , the balance being Fe and impurities is heated. A method for producing a bi-directional electrical steel sheet, comprising a step of hot rolling, cold rolling, and annealing in a two-phase region under reduced pressure by interposing an annealing separator between steel plates, Is heated in the middle to the (α + γ) two-phase region of 750 ° C. or higher, and subjected to intermediate annealing in which the cooling time from immediately above the A1 point to 500 ° C. during cooling is 2 minutes or less. A method for producing a bi-directional electrical steel sheet. The steel containing, by mass%, C: 0.02 to 0.20%, Si: 2.4 to 4.0%, Mn: 0.20 to 2.0% , the balance being Fe and impurities is heated. A method for producing a bi-directional electrical steel sheet, comprising a step of hot rolling, cold rolling, and annealing in a two-phase region under reduced pressure by interposing an annealing separator between steel plates, Satisfies the relationship that the ratio (D / t) between the diameter D (mm) of the work roll and the steel sheet thickness t (mm) during cold rolling is D / t ≧ 80 or D / t ≧ 100 / t. In the middle of this, heating is performed in the (α + γ) two-phase region of 750 ° C. or higher, and intermediate annealing is performed in which the cooling time from immediately above the A1 point to 500 ° C. during cooling is 2 minutes or less. A method for producing a bi-directional electrical steel sheet, which is characterized.   The method for producing a bidirectional steel sheet according to any one of claims 1 to 3, wherein a rolling reduction per pass of cold rolling is 25% or less.   The method for producing a bi-directional electrical steel sheet according to any one of claims 1 to 4, wherein the rolling is performed so that a friction coefficient μ between the work roll and the steel sheet during cold rolling is 0.10 or less. . Soaking temperature of intermediate annealing (T, ° C), soaking time (s, sec), average heating rate from 750 ° C to soaking temperature (Vu, ° C / sec) and average from soaking temperature to 750 ° C The intermediate temperature annealing is performed by selecting these conditions so that the soaking coefficient G calculated by the following formula from the cooling rate (Vd, ° C./sec) is 4500 (° C. * sec) or more. The manufacturing method of the bidirectional electrical steel sheet in any one of 2-5.

Images