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JP2004043921A - Rare-earth-containing alloy flake, its manufacturing process, rare-earth sintered magnet, alloy powder for this, bond magnet and alloy powder for this - Google Patents

Rare-earth-containing alloy flake, its manufacturing process, rare-earth sintered magnet, alloy powder for this, bond magnet and alloy powder for this Download PDF

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JP2004043921A
JP2004043921A JP2002205214A JP2002205214A JP2004043921A JP 2004043921 A JP2004043921 A JP 2004043921A JP 2002205214 A JP2002205214 A JP 2002205214A JP 2002205214 A JP2002205214 A JP 2002205214A JP 2004043921 A JP2004043921 A JP 2004043921A
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alloy
rare earth
rich phase
fine
casting
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JP2004043921A5 (en
Inventor
Shiro Sasaki
佐々木 史郎
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Resonac Holdings Corp
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Showa Denko KK
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Priority to JP2002205214A priority Critical patent/JP2004043921A/en
Priority to AU2002358316A priority patent/AU2002358316A1/en
Priority to US10/498,932 priority patent/US7442262B2/en
Priority to PCT/JP2002/013231 priority patent/WO2003052778A1/en
Priority to CNB028050975A priority patent/CN1306527C/en
Publication of JP2004043921A publication Critical patent/JP2004043921A/en
Publication of JP2004043921A5 publication Critical patent/JP2004043921A5/ja
Priority to US11/826,114 priority patent/US7571757B2/en
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  • Manufacture Of Metal Powder And Suspensions Thereof (AREA)
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Abstract

<P>PROBLEM TO BE SOLVED: To manufacture an alloy mass having a metallic structure with an excellent homogeneity by inhibiting formation of regions containing fine R-rich phase area in a cast R-T-B alloy mass, and to provide a rare-earth magnet showing excellent magnetic characteristics. <P>SOLUTION: In a process for manufacturing the rare earth-containing alloy flake through strip casting, a roll kneader for casting with a casting surface on which two or more mutually crossing linear irregularities are formed and which has a surface roughness of 3-30 μm in terms of ten-point average roughness (Rz), is used. <P>COPYRIGHT: (C)2004,JPO

Description

【0001】
【発明の属する技術分野】
本発明は、希土類含有合金薄片およびその製造方法に係り、特にストリップキャスト法により製造される希土類含有合金薄片およびその製造方法に関する。
【0002】
【従来の技術】
近年、希土類磁石用合金、特にNd−Fe−B系合金がその高特性から急激に生産量を伸ばしており、HD(ハードディスク)用、MRI(磁気共鳴映像法)用あるいは、各種モーター用等に使用されている。通常は、Ndの一部をPr、Dy等の他の希土類元素で置換したものや、Feの一部をCo、Ni等の他の遷移金属で置換したものが一般的であり、Nd−Fe−B系合金を含め、R−T−B系合金と総称されている。ここで、RはYを含む希土類元素のうち少なくとも1種である。また、TはFeを必須とする遷移金属であり、Feの一部をCoあるいはNiで置換することができ、添加元素としてCu、Al、Ti、V、Cr、Ga、Mn、Nb、Ta、Mo、W、Ca、Sn、Zr、Hfなどを1種または複数組み合わせて添加してもよい。Bは硼素であり、一部をCまたはNで置換できる。
【0003】
R−T−B系合金は、磁化作用に寄与する強磁性相であるR14B相からなる結晶を主相とし、非磁性で希土類元素の濃縮した低融点のR−リッチ相が共存する合金で、活性な金属であることから一般に真空又は不活性ガス中で溶解や鋳造が行われる。また、鋳造されたR−T−B系合金塊から粉末冶金法によって焼結磁石を作製するには、合金塊を3μm(FSSS:フィッシャーサブシーブサイザーでの測定)程度に粉砕して合金粉末にした後、磁場中でプレス成形し、焼結炉で約1000〜1100℃の高温にて焼結し、その後必要に応じ熱処理、機械加工し、さらに耐食性を向上するためにメッキを施し、焼結磁石とするのが普通である。
【0004】
R−T−B系合金からなる焼結磁石において、R−リッチ相は、以下のような重要な役割を担っている。
1)融点が低く、焼結時に液相となり、磁石の高密度化、従って磁化の向上に寄与する。
2)粒界の凹凸を無くし、逆磁区のニュークリエーションサイトを減少させ保磁力を高める。
3)主相を磁気的に絶縁し保磁力を増加する。
従って、成形した磁石中のR−リッチ相の分散状態が悪いと局部的な焼結不良、磁性の低下をまねくため、成形した磁石中にR−リッチ相が均一に分散していることが重要となる。ここでR―リッチ相の分布は、鋳造された際のR−T−B系合金塊の組織に大きく影響される。
【0005】
また、R−T−B系合金の鋳造において生じるもう一つの問題は、鋳造された合金塊中にα―Feが生成することである。α―Feは、合金塊を粉砕する際の粉砕効率の悪化をもたらし、また焼結後も磁石中に残存すれば、磁石の磁気特性の低下をもたらす。そこで従来の合金塊では、必要に応じ高温で長時間にわたる均質化処理を行い、α―Feの消去を行っていた。
【0006】
この鋳造されたR−T−B系合金塊中にα−Feが生成する問題を解決するため、より速い冷却速度で合金塊を鋳造する方法として、ストリップキャスト法(SC法と略す。)が開発され実際の工程に使用されている。
SC法は内部が水冷された銅製の鋳造用回転ロール上にR−T−B系合金の溶湯を流し、0.1〜1mm程度の薄片を鋳造することにより、合金を急冷凝固させるものであり、α‐Feの析出を抑制することができる。さらに合金塊の結晶組織が微細化するため、R−リッチ相が微細に分散した組織を有する合金を生成することが可能となる。このように、SC法で鋳造されたR−T−B系合金は、内部のR−リッチ相が微細に分散しているため、粉砕、焼結後の磁石中のR−リッチ相の分散性も良好となり、磁石の磁気特性の向上に成功している。(特開平5−222488号公報、特開平5−295490号公報)
【0007】
またSC法により鋳造されたR−T−B系合金塊は、組織の均質性も優れている。組織の均質性は、結晶粒径やRリッチ相の分散状態で比較することが出来る。SC法で作製した合金薄片では、合金薄片の鋳造用ロール側(以降、鋳型面側とする)にチル晶が発生することもあるが、全体として急冷凝固でもたらされる適度に微細で均質な組織を得ることが出来る。
【0008】
以上のように、SC法で鋳造したR−T−B系合金は、Rリッチ相が微細に分散し、α−Feの生成も抑制されているため、焼結磁石を作製する場合には、最終的な磁石中のRリッチ相の均質性が高まり、またα−Feに起因する粉砕、磁性への弊害を防止することができる。このように、SC法で鋳造したR−T−B系合金塊は、焼結磁石を作製するため優れた組織を有している。しかし、磁石の特性が向上するにつれて、ますます原料合金塊の組織に均質性の向上が求められるようになってきている。
【0009】
そのため、例えば特開平10−317110号公報には、鋳造されたR−T−B系合金の鋳型面側のチル晶の面積比率を5%以下にすることで、磁石特性の良好な焼結磁石を作製している技術が開示されている。チル晶部は粉砕工程で粒径1μm以下の微細粉末となるため、合金粉末の粒度分布を乱し、磁性を悪化させると考えられている。
【0010】
先に本発明者らは、鋳造されたR−T−B系合金塊の組織と、水素解砕や微粉砕の際の挙動との関係を研究した結果、焼結磁石用の合金粉末の粒度を均一に制御するためには、合金塊の結晶粒径よりもRリッチ相の分散状態を制御することが重要であることを見出した(特願2001−383989号)。そして、合金塊中のチル晶の体積率は現実には数%以下であり、チル晶による弊害よりも、合金塊中の鋳型面側に生成されるRリッチ相の分散状態が極端に細かな領域(微細Rリッチ相領域)の方が、磁石用粉末の粒度を制御するためには影響が大きいことを見出した。すなわち、合金塊の組成や製造条件によりR−T−B系合金塊中のチル晶を少なくした場合でも、微細Rリッチ相領域の体積率が50%を超える場合もあること、そしてこの微細Rリッチ相領域が磁石用合金粉末の粒度分布を乱すことを確認し、微細Rリッチ相領域を減少させることが磁石特性を向上させるために必要であることを確認した。
【0011】
そして、SC法における鋳造条件、特に鋳造用回転ロールの表面状態を変更し、R−T−B系合金薄片中の微細Rリッチ相領域が生成する体積率を比較した。すると合金薄片の鋳型面側表面の表面粗さと微細Rリッチ相領域が生成する体積率に関係があることを見出し、微細Rリッチ相領域が20%以下である組織の均質性に優れた合金薄片の製造を可能とした。
【0012】
【発明が解決しようとする課題】
上記の特願2001−383989号に示す方法で、微細Rリッチ相領域を減少させ、組織を均質化することは、ある程度可能である。しかし、鋳造ロールの表面状態以外にも組織に影響をもたらす様々な要因が存在するため、実際のR−T−B系合金の製造ではそれらの要因をすべて管理することは困難であり、合金の一部に微細Rリッチ相領域が生成することがある。そこで本発明は、鋳造されたR−T−B系合金塊中での微細Rリッチ相領域の生成を従来の方法より効果的に抑制し、均質性に優れた組織を有する合金塊を製造する方法を提供する。そして、磁石中のRリッチ相の分布の均質性をさらに高め、磁石特性の優れた希土類磁石を提供することを目的とする。
【0013】
【課題を解決するための手段】
本発明者らは、SC法における鋳造用回転ロールの表面状態を変更し、R−T−B系合金薄片中の微細Rリッチ相領域が生成する体積率を比較した。すると合金薄片の鋳型面側表面の表面粗さに加え、その凹凸の形状と微細Rリッチ相領域が生成する体積率に関係があることを見出した。本発明は、本発明者らが上記の知見に基づき為したものである。
【0014】
すなわち本発明は、
(12)厚さが0.1mm以上0.5mm以下であり、少なくとも片側の表面に互いに交差するように形成された複数の線状の凹凸を有し、線状の凹凸を有する面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下であることを特徴とする希土類含有合金薄片。
(2)希土類磁石用原料として使用されるR−T−B系(但し、RはYを含む希土類元素のうち少なくとも1種、TはFeを必須とする遷移金属、Bは硼素である。)合金からなることを特徴とする上記(1)に記載の希土類含有合金薄片。
(3)合金中の微細Rリッチ相領域の体積率が20%以下であることを特徴とする上記(2)に記載の希土類含有合金薄片。
(4)ストリップキャスト法(SC法)による希土類含有合金薄片の製造方法において、鋳造面に互いに交差するように複数の線状の凹凸が形成され、該鋳造面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下である鋳造用回転ロールを用いることを特徴とする希土類含有合金薄片の製造方法。
(5)希土類含有合金薄片が、厚さが0.1mm以上0.5mm以下であり、少なくとも片側の表面に互いに交差するように形成された複数の線状の凹凸を有し、線状の凹凸を有する面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下であることを特徴とする上記(4)に記載の希土類含有合金薄片の製造方法。
(6)希土類含有合金薄片が、希土類磁石用原料として使用されるR−T−B系(但し、RはYを含む希土類元素のうち少なくとも1種、TはFeを必須とする遷移金属、Bは硼素である。)合金からなることを特徴とする上記(4)または(5)に記載の希土類含有合金薄片の製造方法。
(7)合金中の微細Rリッチ相領域の体積率が20%以下であることを特徴とする上記(6)に記載の希土類含有合金薄片の製造方法。
(8)上記(2)又は(3)に記載の希土類含有合金薄片に水素解砕工程を施し、その後にジェットミル粉砕することで作製される希土類焼結磁石用合金粉末。
(9)上記(8)に記載の希土類焼結磁石用合金粉末から粉末冶金法で製造される希土類焼結磁石。
(10)上記(2)又は(3)に記載の希土類含有合金薄片を用いて、HDDR法で製造したボンド磁石用合金粉末。
(11)上記(10)に記載のボンド磁石用合金粉末を用いて作製されるボンド磁石。
である。
【0015】
【発明の実施の形態】
従来のSC法により鋳造されたNd−Fe−B系合金(Nd31.5質量%)の薄片の断面をSEM(走査電子顕微鏡)にて観察した時の反射電子像を図1に示す。図1で左側が鋳型面側、右側が自由面側である。なお、この合金薄片の鋳型面側表面の表面粗さは十点平均粗さ(Rz)で3.4μmであり、該表面には直線状の凹凸がある方向にほぼ平行に付いていた。
図1で白い部分が、Nd−リッチ相(RがNdになっているためR−リッチ相をNd−リッチ相と呼ぶ。)で、合金薄片の中央部から自由面側(鋳造面側と反対側の表面)では、厚さ方向にラメラー状に伸びるか、ラメラーが分断したような方向性を持った形の小さなプールを形成している。しかし、鋳型面側にはNd−リッチ相が他の部位よりも極端に微細な粒状で、かつランダムに存在する領域が生成しており、これを本発明者らは微細Rリッチ相領域(Rの主成分がNdの際は微細Ndリッチ相領域とも呼ぶ)と名づけ、特に区別することとした。この微細Rリッチ相領域は通常鋳型面側から始まり、中央方向へ広がっている。これに対し中央部から自由面側にかけての微細Rリッチ相領域が存在しない部分を、ここでは正常部と呼ぶこととする。
【0016】
焼結磁石作製時のR−T−B系合金薄片の水素解砕工程において、水素はRリッチ相から吸収され、膨張し脆い水素化物となる。したがって、水素解砕では、合金中にRリッチ相に沿った、或いはRリッチ相を起点とした微細なクラックが導入される。その後の微粉砕工程で、水素解砕で生成した多量の微細クラックをきっかけに合金が壊れるため、合金中のRリッチ相の分散が細かいほど微粉砕後の粒度は細かくなる傾向がある。したがって、微細Rリッチ相領域は、正常部よりも細かく割れる傾向が強く、例えば正常部から製造された合金粉末では、平均粒度がFSSS(フィッシャー サブ シーブ サイザー)での測定で3μm程度であるのに対して、微細Rリッチ相領域から製造された合金粉末では、1μm以下の微粉を含む割合が高いため、微粉砕後の粒度分布が広くなることになる。
【0017】
R−T−B系合金中のRリッチ相の分散状態は、鋳造時における溶湯が凝固した後の冷却速度の制御、或いは熱処理によって制御可能であることは特開平09−170055号公報、或いは特開平10−36949号公報に記載されている。しかし、凝固後の冷却速度、或いは熱処理による微細Rリッチ相領域内部のRリッチ相の変化の挙動は、正常部と異なり制御が困難であり、Rリッチ相の分散が粗くなりにくく、微細なままである。
【0018】
微細Rリッチ相領域の体積率は次のような方法で測定可能である。図3は図1と同じ視野の反射電子線像であるが、微細Rリッチ相領域と正常部の境界に線を引いたものである。両領域の境界は、Rリッチ相の分散状態から容易に判断できるため、画像解析装置を用いてその視野の微細Rリッチ相領域の面積率を計算することが出来る。断面での面積率は、合金中での体積率に対応する。なお、微細Rリッチ相領域の体積率の測定において、同時に鋳造された合金薄片であっても、微細Rリッチ相領域の量の変化は、薄片間同士、また同じ薄片内でも大きい。そのため、50〜100倍程度の低倍率で観察視野を広げた上で、5〜10枚程度の薄片を測定しその平均を取ることで、その合金全体の微細Rリッチ相領域の体積率を計算することが出来る。
【0019】
本発明のR−T−B系合金薄片(Nd31.5質量%)の断面の反射電子線像を図2に示す。図2で左側が鋳型面側、右側が自由面側である。本発明の合金薄片の特徴は、ストリップキャスト法で製造された薄片において、鋳型面側の表面粗さを制御すると同時に、該表面の線状の凹凸を互いに交差するように形成することによって、微細Rリッチ相領域の生成が抑制されていることである。図2に示す合金薄片の鋳型面側の表面粗さは3.2μmと、図1の合金薄片とほぼ同じである。しかし、本発明の合金薄片では、鋳型面側に微細Rリッチ相領域は存在せず、鋳型面から自由面に渡ってRリッチ相の分散状態が極めて均質である。
【0020】
ストリップキャスト法で製造された合金薄片の鋳型面側表面の表面粗さと線状の凹凸の状態、微細Rリッチ相領域の関係は以下のように説明できる。
合金薄片の鋳型面側表面が平滑であるためには、鋳造用回転ロール表面が平滑で、合金溶湯との濡れ性が良好である必要がある。このような状態では、溶湯から鋳型への熱伝達が極めて良好(熱伝達係数が大きい)であり、合金の鋳型面側が過度に急冷される。微細Rリッチ相領域は、鋳型と溶湯の熱伝達係数が大きく合金の鋳型面側が過度に急冷される場合に生成される傾向が強いと考えられる。
【0021】
一方、鋳造用回転ロールの表面に細かな凸凹を形成すると、合金溶湯の粘性のため、溶湯は鋳造用回転ロール表面の細かな凸凹に完全には入り込めず、未接触の部分を生じ、熱伝達係数が低下する。その結果、合金の鋳型面側が過度に急冷されることがなくなり、微細Rリッチ相領域の生成が抑制できると考えられる。ここで鋳造用回転ロール表面の表面粗さを大きくすると、合金薄片の鋳型面側に多少なりともその凸凹が転写されるため、合金薄片の鋳型面側表面の表面粗さも当然大きくなる。鋳型面側表面が適当な表面粗さを有する合金薄片で、Rリッチ相の生成が抑制される原因は、上記のように溶湯が凝固する時の過度の熱伝達が抑制されているためと推定される。
【0022】
ここで、さらに凹凸の形状に注目すると、それが互いに交差しない線状であった場合、溶湯とロールの接触部、非接触部のそれぞれが、線状の凹凸に沿って連続的に存在する傾向がある。したがって、内部組織もその凹凸に沿って連続性を有する傾向にある。その場合、ある線状の凹凸部に何らかの原因で、微細Rリッチ相領域を生成すると、その線状の凹凸部の全体に渡って、微細Rリッチ相領域が生成される危険性が高くなる。
【0023】
しかし、線状の凹凸が互いに交差していると、表面の凹凸が分断されるだけでなく、その交差部で合金の内部組織の連続性も途切れることとなる。さらにこの交差部では、線状の凸部も必ず凹部で切られることになる。この凸部は、溶湯と鋳造ロール表面との接触が良好となるため熱伝達係数が大きくなり、急冷凝固による微細Rリッチ相領域が生成されやすいと考えられる。したがって、微細Rリッチ相が生成されたとしても、その拡大を防止する効果がある。
【0024】
本発明のストリップキャスト法による希土類含有合金薄片の製造方法では、鋳造面に互いに交差するように複数の線状の凹凸が形成され、該鋳造面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下である鋳造用回転ロールを用いる。その結果、少なくとも片側の表面に互いに交差するように形成された複数の線状の凹凸を有し、線状の凹凸を有する面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下である希土類含有合金薄片を製造することが出来る。本発明によれば、従来の方法よりも表面粗さが小さくても、微細Rリッチ相の生成を抑制することができ、組織の均質化の効果を得ることが出来る。また、鋳造用回転ロールの表面粗さを小さくすると、ロール表面を調整する際の切削量を削減出来るため、鋳造用回転ロールの寿命増加を可能とする。また、表面粗さの相違による影響を受けにくいため、ロール表面状態の管理基準を簡素化することも可能である。
【0025】
従来のSC法でも図2に示すような均質な組織を有する合金薄片はある程度含まれていたが、図1に示すような微細Rリッチ相領域を多量に含んだ薄片も同時に生成されてしまうため、結果として合金全体での組織の均質性に問題を生じていた。このような従来のSC法で作製した合金組織のばらつきは、微妙な鋳造用回転ロールの表面状態、溶湯の供給状態、雰囲気など、ロール表面と溶湯との接触状態の違いに起因するものと考えられる。鋳造用回転ロール表面に形成した凸凹は、溶湯が凝固する時の過度の熱伝達を抑制し、微細Rリッチ相領域の生成を再現良く抑制することができる。
さらに本発明では、鋳造用回転ロールの表面の凹凸の形状を互いに交差する線状とすることによって、微細Rリッチ相領域の生成を抑制する効果を高めたため、比較的小さな表面粗さでも効果を得ることができた。その結果、図2に示すような均質な組織を有する合金薄片の収率をより一層大きくすることができた。
【0026】
さらに本発明の詳細を説明する。
(1)ストリップキャスト法(SC法)
本発明は希土類含有合金薄片に関するものである。ここでは、R−T−B系合金のストリップキャスト法による鋳造について説明する。
図4にストリップキャスト法による鋳造に用いる装置の模式図を示す。通常、R−T−B系合金は、その活性な性質のため真空または不活性ガス雰囲気中で、耐火物ルツボ1を用いて溶解される。溶解された合金の溶湯は1350〜1500℃で所定の時間保持された後、必要に応じて整流機構、スラグ除去機構を設けたタンディッシュ2を介して、内部を水冷された鋳造用回転ロール3に供給される。溶湯の供給速度と回転ロールの回転数は、求める合金の厚さに応じて適当に制御させる。一般に回転ロールの回転数は、周速度にして1〜3m/s程度である。鋳造用回転ロールの材質は、熱伝導性がよく入手が容易である点から銅、或いは銅合金が適当である。回転ロールの材質やロールの表面状態によっては、鋳造用回転ロールの表面にメタルが付着しやすいため、必要に応じて清掃装置を設置すると、鋳造されるR−T−B系合金の品質が安定する。回転ロール上で凝固した合金4はタンディッシュの反対側でロールから離脱し、捕集コンテナ5で回収される。この捕集コンテナに加熱、冷却機構を設けることで正常部のRリッチ相の組織の状態を制御できる。
【0027】
本発明の希土類含有合金薄片の厚さは、0.1mm以上0.5mm以下とするのが好ましい。合金薄片の厚さが0.1mmより薄いと凝固速度が過度に増加し、結晶粒径が細かくなりすぎ、磁石化工程での微粉砕粒度近くになるため、磁石の配向率、磁化の低下を招くという問題がある。また合金薄片の厚さが0.5mmより厚いと凝固速度低下によるNd−rich相の分散性の低下、α‐Feの析出などの問題を招く。
【0028】
(2)鋳造用回転ロールの鋳造面の表面粗さ
本発明においては、ストリップキャスト法でR−T−B系合金を鋳造する場合、鋳造用回転ロールの鋳造面の表面粗さを、十点平均粗さ(Rz)で3μm以上30μm以下とする。
ここで表面粗さとは、JIS B 0601「表面粗さの定義と表示」に示される条件で測定したもので、十点平均粗さ(Rz)もその中に定義されている。具体的にはまず、測定面に直角な平面で切断したときの切り口(断面曲線)から、所定の波長より長い表面うねり成分を位相補償型高域フィルタ等で除去した曲線(粗さ曲線)を求める。その粗さ曲線から、その平均線の方向に基準長さだけ抜き取り、この抜き取り部分の平均線から、最も高い山頂から5番目までの山頂の標高(Yp)の絶対値の平均値と、最も低い谷底から5番目までの谷底の標高(Yv)の絶対値の平均値との和を十点平均粗さ(Rz)と呼ぶ。基準長さ等の測定パラメータは、表面粗さに対して標準値が上記JISで指定されている。
合金薄片の鋳型面側の表面粗さは、変動が大きい場合もあり、少なくとも5枚の薄片について測定し、その平均値を使用すべきである。
【0029】
(3)鋳造用回転ロールの鋳造面の凹凸の形状
本発明では、鋳造面の表面粗さが主として該表面に形成された複数の線状の凹凸によってもたらされるものであり、該線状の凹凸は互いに交差するように表面に形成されている。
凹凸が線状であった場合、溶湯とロールの接触部、非接触部のそれぞれが、線状の凹凸に沿って連続的に存在する傾向がある。したがって、内部組織もその凹凸に沿って連続性を有する傾向にある。その場合、ある線状の凹凸部に何らかの原因で、微細Rリッチ相領域を生成すると、その線状部の全体に渡って、微細Rリッチ相領域が生成される危険性が高くなる。
しかし、線状の凹凸が互いに交差していると、表面の凹凸が分断されるだけでなく、その交差部で内部組織の連続性も途切れることとなり、微細Rリッチ相が生成されたとしても、その拡大を防止する効果がある。
【0030】
本発明では、互いに交差するように存在する線状の凹凸の効果で、表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下と比較的小さくても、組織を均質化する効果を得ることが出来る。
しかし、表面粗さが3μm以下では凸凹の効果が得られず、鋳造用回転ロールの表面と合金溶湯との接触が良好なため、熱伝達係数が大きくなる。その結果、合金中に微細Rリッチ相領域を生成しやすくなる。
【0031】
鋳造用回転ロールの表面粗さが30μmを超えると、凝固した合金薄片がロール表面に噛み込んで、剥がれ難くなり、タンディッシュを破壊するなどのトラブルの原因となることがある。そのため鋳造用回転ロールの表面粗さは、30μm以下する。
【0032】
(4)希土類含有合金薄片の表面粗さと凹凸の形状
本発明においては、希土類含有合金薄片の少なくとも片面の表面粗さが、十点平均粗さ(Rz)で3μm以上30μm以下であることを特徴とし、この表面粗さが主として該表面に形成された複数の線状の凹凸によってもたらされるものであり、該線状の凹凸は互いに交差するように表面に形成されていることを特徴とする。
表面に上記の粗さの凸凹が形成される面は、ストリップキャスト法で鋳造する際に凝固が始まる鋳型面側表面であり、回転ロールの表面の凸凹が反映された表面となる。上記した通り、この表面の表面粗さが3μm以下では、微細Rリッチ相領域が生成する体積率が大きくなり、合金中のRリッチ相の分散状態の不均一をもたらす。その結果、焼結磁石の製造工程で微粉砕後の合金粉末の粒度分布を広くし、磁石の特性を悪化するため好ましくない。一方、30μm以上では合金の鋳造の過程で問題を生じやすい。本発明において合金薄片の片面の表面粗さは、3μm以上30μm以下とする。
【0033】
(5)合金中の微細Rリッチ相領域の体積率
本発明では、R−T−B系合金中の微細Rリッチ相領域の体積率は20%以下となる。その結果、焼結磁石の製造工程で、微粉砕後の合金粉末の粒度分布が狭く揃ったものになるため、特性にバラツキのない均質な焼結磁石を得ることができる。
【0034】
(6)希土類焼結磁石用合金粉末および希土類焼結磁石の製造方法
本発明により鋳造した磁石用原料のR−T−B系合金からなる希土類含有合金薄片からは、粉砕、成型、焼結の工程を経て、高特性の異方性焼結磁石を製造することができる。
【0035】
合金薄片の粉砕は、通常、水素解砕、微粉砕の順で行なわれ、3μm(FSSS)程度の合金粉末が作製される。ここで、水素解砕は、前工程の水素吸蔵工程と後工程の脱水素工程に分けられる。水素吸蔵工程では、266hPa〜0.3MPaの圧力の水素ガス雰囲気で、主に合金薄片のR−リッチ相に水素を吸蔵させ、この時に生成されるR−水素化物によりR−リッチ相が体積膨張することを利用して、合金薄片自体を微細に割るかあるいは無数の微細な割れ目を生じさせる。この水素吸蔵は常温〜600℃程度の範囲で実施されるが、R−リッチ相の体積膨張を大きくして効率良く割るためには、水素ガス雰囲気の圧力を高くすると共に、常温〜100℃程度の範囲で実施することが好ましい。好ましい処理時間は1時間以上である。この水素吸蔵工程により生成したR−水素化物は大気中では不安定であり酸化され易いため、水素吸蔵処理の後、200〜600℃程度で1.33hPa以下の真空中に合金薄片を保持する脱水素処理を行なうことが好ましい。この処理により、大気中で安定なR−水素化物に変化させることができる。脱水素処理の好ましい処理時間は30分以上である。水素吸蔵後から焼結までの各工程で酸化防止のための雰囲気管理がなされている場合は、脱水素処理を省くこともできる。
【0036】
本発明のストリップキャスト法により製造されたR−T−B系合金薄片は、Rリッチ相が均一に分散していることが特徴である。好ましいRリッチ相の間隔の平均値は、磁石の製造工程での粉砕粒度に依存するが、一般に3μmから8μmである。水素解砕では、Rリッチ相に沿って、或いはRリッチ相を起点にして合金内にクラックが導入される。したがって、水素解砕してから微粉砕することで、合金中に均一かつ微細に分散したRリッチ相の効果を最大限に引き出すことが可能であり、非常に粒度分布の狭い合金粉末を効率良く生産することが可能である。この水素解砕の工程を行わずに焼結磁石を作製した場合、作製された焼結磁石の特性は劣ったものとなる。(M.Sagawa et al. Proceeding of the 5th international conference on Advanced materials,Beijing China(1999))
【0037】
微粉砕とは、R−T−B系合金薄片を3μm(FSSS)程度まで粉砕することである。微粉砕のための粉砕装置としては、生産性が良く、狭い粒度分布を得られることから、ジェットミル装置が最適である。本発明の微細Rリッチ相領域の少ない合金薄片を利用すれば、粒度分布が狭い合金粉末を高効率で安定性良く作製することができる。
微粉砕を行う際の雰囲気は、アルゴンガスや窒素ガスなどの不活性ガス雰囲気とする。これらの不活性ガス中に2質量%以下、好ましくは1質量%以下の酸素を混入させてもよい。このことにより粉砕効率が向上するとともに、粉砕後の合金粉末の酸素濃度を1000〜10000ppmとすることができ、合金粉末を適度に安定化させることができる。また同時に、磁石を燒結する際の結晶粒の異常成長を抑制することもできる。
【0038】
上記の合金粉末を磁場中で成型する場合、合金粉末と金型内壁との摩擦を低減し、また粉末どうしの摩擦も低減させて配向性を向上させるため、合金粉末にはステアリン酸亜鉛等の潤滑剤を添加することが好ましい。好ましい添加量は0.01〜1質量%である。潤滑材の添加は微粉砕前でも後でもよいが、磁場中成形前に、アルゴンガスや窒素ガスなどの不活性ガス雰囲気中でV型ブレンダー等を用いて十分に混合することが好ましい。
【0039】
3μm(FSSS)程度まで粉砕されたR−T−B系合金粉末は、磁場中成型機でプレス成型される。金型は、キャビティ内の磁界方向を考慮して、磁性材と非磁性材を組み合わせて作製される。成型圧力は0.5〜2t/cmが好ましい。成型時のキャビティ内の磁界は5〜20kOeが好ましい。また、成型時の雰囲気はアルゴンガスや窒素ガスなどの不活性ガス雰囲気が好ましいが、上述の耐酸化処理した粉体の場合、大気中でも可能である。
また成形は、冷間静水圧成形(CIP:Cold Isostatic Press)或いはゴム型を利用した擬似静水圧プレス(RIP:Rubber Isostatic Press)でも可能である。CIPやRIPでは、静水圧的に圧縮されるため、成形時の配向の乱れが少なく、金型成形よりも配向率の増加が可能であり、最大磁気エネルギー積を増加することができる。
【0040】
成型体の焼結は、1000〜1100℃で行なわれる。焼結の雰囲気としては、アルゴンガス雰囲気または1.33×10−2hPa以下の真空雰囲気が好ましい。焼結温度での保持時間は1時間以上が好ましい。また焼結の際には、焼結温度に到達する前に、成型体中の潤滑剤と合金粉末に含まれる水素はできるだけ除去しておく必要がある。潤滑剤の好ましい除去条件は、1.33×10−2hPa以下の真空中または減圧したArフロー雰囲気中で、300〜500℃で30分以上保持することである。また、水素の好ましい除去条件は、1.33×10−2hPa以下の真空中で、700〜900℃で30分以上保持することである。
【0041】
焼結が終了した後、焼結磁石の保磁力向上のため、必要に応じて500〜650℃で熱処理することができる。この場合の好ましい雰囲気は、アルゴンガス雰囲気または真空雰囲気であり、好ましい保持時間は30分以上である。
【0042】
また、本発明で作製した微細Rリッチ領域の生成を抑制した希土類磁石用R−T−B系合金薄片は、焼結磁石以外に、ボンド磁石の作製のためにも好適に用いることができる。以下に、本発明の希土類磁石用合金薄片からボンド磁石を作製する場合について説明する。
【0043】
本発明のR−T−B系合金薄片は、まず必要に応じて熱処理される。熱処理の目的は、合金中のα‐Feの除去と結晶粒の粗大化である。ボンド磁石のための合金粉末の作製には、HDDR(Hydrogenation Disproportionation Desorption Recombination)処理を行うが、合金中に存在するα‐FeはHDDR処理工程では消去させることができず、磁性を低下させる原因となる。そのため、α−FeはHDDR処理を行う前に消去しておく必要がある。
【0044】
また、ボンド磁石用の合金粉末の平均粒径は50〜300μmと焼結磁石用の合金粉末と比較すると非常に大きい。HDDR法では、元の合金の結晶方位と、再結合したサブミクロンの結晶粒の方位がある一定の分布を持って一致する。そのため、原料の合金薄片中にある二つ以上の結晶方位の異なる結晶粒が、一つのボンド磁石用合金粉末に含まれてしまうと、合金粉末中に結晶方位が大きく異なる領域を含むこととなり、磁石の配向率が低下し、最大磁気エネルギー積が低下する。これを避けるためには、合金薄片中の結晶粒径は、大きい方が都合が良い。ストリップキャスト法のような急冷凝固法で鋳造した合金では、結晶粒径が比較的小さくなる傾向があるため、熱処理による結晶粒の粗大化はボンド磁石の特性の向上に有効である。
【0045】
HDDR法によるボンド磁石用合金粉末の製造方法については、多くの報告がある(例えば、T.Takeshita et al,Proc.10th Int. Workshop on RE magnets and their application, Kyoto, Vol.1 p551(1989))。HDDR法による合金粉末の作製は、以下のように行われる。
【0046】
原料のR−T−B系合金薄片を水素雰囲気中で加熱すると、700℃から850℃程度で磁性相のR14B相がα‐Fe、RH、FeBの3相に分解する。次いで同程度の温度で、不活性ガス雰囲気、或いは真空雰囲気に切り替えて水素を除去すると、分解していた相がサブミクロン程度の結晶粒径を有するR14B相に再結合する。この際、合金の組成や処理条件を適当に制御すると、再結合した各R14B相の磁化容易軸(R14B相C軸)は、分解前の原料合金中のR14B相のC軸とほぼ平行となり、各微細結晶粒の磁化容易軸方向が揃った異方性磁石粉とすることができる。
【0047】
HDDR処理を施した合金は、50〜300μm程度に粉砕し合金粉末とした後、樹脂と混合して圧縮成形、射出成形などを施しボンド磁石とすることできる。
【0048】
微細Rリッチ相領域は上記した水素解砕処理同様に、HDDR処理の際にも微粉化する傾向が強い。HDDR法による磁粉の特性は、粒度が小さくなるとともに低下する。そのため、本発明の微細Rリッチ相の生成を抑制したR−T−B系合金は、HDDR処理でのボンド磁石用磁粉の作製に好適に用いることができる。
【0049】
最近、SC法による希土類合金の製造において、鋳造用回転ロールの外周面の表面粗さのSm/RaとSmをある範囲に制御して、合金組織の均質性を改善した報告がなされた(特開2002−59245号公報、特開平9−1296号公報)。しかし、それらはストリップの幅方向での組織変化の抑制を目的とし、ストリップ端部での冷却速度の低下を防止しようとするものである。また、表面粗さをもたらす凹凸の形状に付いては、特に定めていない。
一方、本発明は合金薄片の厚さ方向、すなわちロール面側から自由面側への組織変化を抑制し、組織の均質化を図るものであり、均質性の尺度として微細Rリッチ相領域を導入し、具体的な体積率を明示したものであり、特開2002−59245号公報或いは特開平9−1296号公報の発明とは、全く異なっている。
【0050】
【実施例】
(実施例1)
合金組成が、Nd:31.5質量%、B:1.00質量%、Co:1.0質量%、Al:0.30質量%、Cu:0.10質量%、残部鉄になるように、金属ネオジウム、フェロボロン、コバルト、アルミニウム、銅、鉄を配合した原料を、アルミナ坩堝を使用して、アルゴンガスで1気圧の雰囲気中で、高周波溶解炉で溶解し、溶湯をストリップキャスト法にて鋳造して、合金薄片を作製した。
鋳造用回転ロールの直径は300mm、材質は純銅で、内部は水冷されており、鋳造面の表面粗さは十点平均粗さ(Rz)で4.0μmに調整した。この際の表面粗さは、主に鋳造面に付与された線状の互いに交差するような無秩序の方向の凹凸によってもたらされた。鋳造時のロールの周速度は1.0m/sで、平均厚さ0.30mmの合金薄片を生成した。
【0051】
得られた合金薄片の鋳型面側表面の表面粗さは、十点平均粗さ(Rz)で4.6μmであり、互いに交差するような線状の凹凸を有していた。合金薄片を10枚埋め込み、研摩した後、走査型電子顕微鏡(SEM)で各合金薄片について反射電子線像(BEI)を倍率100倍で撮影した。撮影した写真を画像解析装置に取り込み測定したところ、微細Rリッチ相領域の体積率は、3%以下であった。
【0052】
(比較例1)
実施例1と同様の組成に原料を配合し、実施例1と同様にして溶解およびSC法による鋳造を実施した。但し、鋳造用回転ロール表面の表面粗さは十点平均粗さ(Rz)で4.0μmであり、ロールの回転方向にほぼ平行な線状の凹凸を有し、それに交差するような凹凸は実質的には存在しなかった。
得られた合金薄片を実施例1と同様に評価した結果、鋳型面側表面の表面粗さは十点平均粗さ(Rz)で4.5μmであり、微細Rリッチ相領域の体積率は、25%であった。
【0053】
(比較例2)
実施例1と同様の組成に原料を配合し、実施例1と同様にして溶解およびSC法による鋳造を実施した。但し、鋳造用回転ロール表面の表面粗さは十点平均粗さ(Rz)で100μmであり、実施例1と同様にその表面粗さは主に互いに交差するような線状の凹凸によってもたらされていた。
すると、鋳造途中でメタルの一部がロールから離脱しないまま1周して、タンディッシュに接触、その前端を破壊した為、鋳造作業を中止した。
【0054】
次に焼結磁石を作製した実施例を説明する。
(実施例2)
実施例1で得られた合金薄片を水素解砕し、ジェットミルで微粉砕した。水素解砕工程の前工程である水素吸蔵工程の条件は、100%水素雰囲気、2気圧で1時間保持とした。水素吸蔵反応開始時の金属片の温度は25℃であった。また後工程である脱水素工程の条件は、0.133hPaの真空中で、500℃で1時間保持とした。この粉末に、ステアリン酸亜鉛粉末を0.07質量%添加し、100%窒素雰囲気中でV型ブレンダーで十分混合した後、ジェットミル装置で微粉砕した。粉砕時の雰囲気は、4000ppmの酸素を混合した窒素雰囲気中とした。その後、再度、100%窒素雰囲気中でV型ブレンダーで十分混合した。得られた粉体の酸素濃度は2500ppmで、粉体の炭素濃度の分析から、粉体に混合されているステアリン酸亜鉛粉末は0.05質量%であると計算された。また、レーザー回折式粒度分布測定機で測定した結果、平均粒度D50は5.00μm、D10は1.98μm、D90は8.51μmであった。
【0055】
次に、得られた粉体を100%窒素雰囲気中で横磁場中成型機でプレス成型した。成型圧は1.2t/cmであり、金型のキャビティ内の磁界は15kOeとした。得られた成型体を、1.33×10−5hPaの真空中、500℃で1時間保持し、次いで1.33×10−5hPaの真空中、800℃で2時間保持した後、さらに1.33×10−5hPaの真空中、1050℃で2時間保持して焼結させた。焼結密度は7.5g/cm以上であり十分な大きさの密度となった。さらに、この焼結体をアルゴン雰囲気中、560℃で1時間熱処理し、焼結磁石を作製した。
【0056】
直流BHカーブトレーサーでこの焼結磁石の磁気特性を測定した結果を表1に示す。また、焼結磁石の原料の微粉の酸素濃度と粒度も表1に示す。
【0057】
(比較例3)
比較例1で得られた合金薄片を、実施例2と同様の方法で粉砕して微粉を得た。さらに実施例2と同様の成型、焼結の工程を経て、焼結磁石を作製した。
【0058】
本比較例3で作製した焼結磁石の磁気特性を、直流BHカーブトレーサーで測定した結果を表1に示す。また、本比較例3の焼結磁石の原料の微粉の酸素濃度と粒度も表1に示す。
【0059】
【表1】

Figure 2004043921
【0060】
表1に示すように、比較例3では実施例2と比較してD10が小さいことから、1μm程度より小さい非常に細かい粉末の割合が大きい事がわかる。このような非常に細かい粒は酸化しやすく、比較例3では実施例2よりも微粉の酸素濃度が若干高くなっている。比較例3の磁石の磁気特性が実施例2と比較して低い原因は、酸素濃度増加と結晶組織の不均質性が主因と考えられる。
【0061】
次にボンド磁石を作製した実施例を説明する。
(実施例3)
合金組成が、Nd28.5%、B:1.00質量%、Co:10.0質量%、Ga:0.5質量%、残部鉄になるように原料を配合し、実施例1と同様の条件でSC法により合金薄片を鋳造した。
得られた合金薄片を実施例1と同様に評価した結果、鋳型面側表面の表面粗さは十点平均粗さ(Rz)で4.3μm、微細Rリッチ相領域の体積率は3%以下であり、α‐Feは含んでいなかった。
【0062】
上記の合金薄片を1気圧の水素中、820℃で1時間保持した後、同温度で真空で1時間保持するHDDR処理を実施した。得られた合金粉を150μm以下にブラウンミルで粉砕し、2.5質量%のエポキシ樹脂を加えて1.5Tの磁場を加えて圧縮成形してボンド磁石を得た。得られたボンド磁石の磁気特性を表1に示す。
【0063】
(比較例4)
実施例3と同様の組成に原料を配合し、比較例1と同様にして溶解およびSC法による鋳造を実施した。得られた合金薄片を実施例1と同様に評価した結果、鋳型面側表面の表面粗さは十点平均粗さ(Rz)で4.8μm、微細Rリッチ相領域の体積率は、30%であった。
【0064】
次いで、本比較例4で得られた合金薄片を用いて、実施例3と同様の方法でボンド磁石を作製した。得られたボンド磁石の磁気特性を表1に示す。
【0065】
表1から本実施例3と比較例4のボンド磁石では、本実施例3の磁気特性が優れていることがわかる。比較例4では、微細Rリッチ領域の体積率が高く、HDDR処理、または粉砕後に50μm以下の比較的細かい粒の量が多いために、磁性が低いものと推定できる。
【0066】
【発明の効果】
本発明のR−T−B系合金薄片は、微細Rリッチ領域の体積率が少なく、合金中のRリッチ相の分散状態の均質性が、従来のSC法により製造した合金薄片よりもさらに良好である。そのため、本発明のR−T−B系合金薄片から製造した焼結磁石やHDDR法によるボンド磁石は、従来のものよりも優れた磁石特性を発現する。
【図面の簡単な説明】
【図1】従来のSC法で製造した微細Rリッチ相を含む希土類磁石用合金薄片の断面組織を示す図である。
【図2】本発明に係る希土類磁石用合金薄片の断面組織を示す図である。
【図3】図1の断面組織における微細Rリッチ領域と正常部との境界に線を引いた図である。
【図4】ストリップキャスト法の鋳造装置の模式図である。
【符号の説明】
1  耐火物ルツボ
2  タンディッシュ
3  鋳造用回転ロール
4  合金
5  捕集コンテナ[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention relates to a rare earth-containing alloy flake and a method for producing the same, and more particularly to a rare earth-containing alloy flake produced by a strip casting method and a method for producing the same.
[0002]
[Prior art]
In recent years, the production of rare earth magnet alloys, especially Nd-Fe-B alloys, has rapidly increased due to their high characteristics, and they are used for HDs (hard disks), MRIs (magnetic resonance imaging), or for various motors. It is used. Normally, Nd-Fe is partially substituted with another rare earth element such as Pr or Dy, or Fe is partially substituted with another transition metal such as Co or Ni. It is collectively referred to as RTB-based alloys, including -B-based alloys. Here, R is at least one of rare earth elements including Y. In addition, T is a transition metal in which Fe is essential, and a part of Fe can be replaced by Co or Ni, and Cu, Al, Ti, V, Cr, Ga, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr, Hf, or the like may be added alone or in combination of two or more. B is boron and can be partially substituted with C or N.
[0003]
The RTB-based alloy is a ferromagnetic phase that contributes to the magnetizing action. 2 T 14 An alloy in which a crystal composed of the B phase is the main phase, and a non-magnetic, rare-earth element-enriched, low-melting R-rich phase coexists. Done. Further, in order to produce a sintered magnet from the cast RTB-based alloy ingot by powder metallurgy, the alloy ingot is pulverized to about 3 μm (FSSS: measured by a Fischer subsieve sizer) to obtain an alloy powder. After that, press forming in a magnetic field, sintering in a sintering furnace at a high temperature of about 1000 to 1100 ° C, heat treatment and machining as needed, and then plating to improve corrosion resistance and sintering Usually it is a magnet.
[0004]
In a sintered magnet made of an RTB-based alloy, the R-rich phase plays an important role as follows.
1) It has a low melting point and becomes a liquid phase at the time of sintering, which contributes to increasing the density of the magnet and thus improving the magnetization.
2) Eliminate irregularities at grain boundaries, reduce nucleation sites in reverse magnetic domains, and increase coercive force.
3) Magnetically insulate the main phase to increase coercive force.
Therefore, if the dispersion state of the R-rich phase in the formed magnet is poor, local sintering failure and a decrease in magnetism may be caused, so it is important that the R-rich phase is uniformly dispersed in the formed magnet. It becomes. Here, the distribution of the R-rich phase is greatly affected by the structure of the R-T-B-based alloy mass at the time of casting.
[0005]
Another problem that occurs in the casting of the RTB-based alloy is that α-Fe is formed in the cast alloy ingot. α-Fe causes deterioration of the pulverizing efficiency when pulverizing the alloy lump, and if it remains in the magnet after sintering, it lowers the magnetic properties of the magnet. Therefore, in the conventional alloy ingot, if necessary, a long time homogenization treatment is performed at a high temperature to eliminate α-Fe.
[0006]
In order to solve the problem that α-Fe is formed in the cast RTB-based alloy ingot, strip casting (abbreviated as SC) is a method of casting the alloy ingot at a higher cooling rate. Developed and used in actual processes.
The SC method is to rapidly solidify an alloy by flowing a melt of an RTB-based alloy on a copper casting rotating roll whose inside is water-cooled and casting a thin piece of about 0.1 to 1 mm. , Α-Fe can be suppressed. Further, since the crystal structure of the alloy lump is refined, an alloy having a structure in which the R-rich phase is finely dispersed can be generated. As described above, in the RTB-based alloy cast by the SC method, since the internal R-rich phase is finely dispersed, the dispersibility of the R-rich phase in the magnet after pulverization and sintering is determined. And the magnetic properties of the magnet were successfully improved. (JP-A-5-222488, JP-A-5-295490)
[0007]
Further, the RTB-based alloy ingot cast by the SC method has excellent structure homogeneity. The homogeneity of the structure can be compared based on the crystal grain size and the dispersion state of the R-rich phase. In alloy flakes produced by the SC method, chill crystals may be generated on the casting roll side (hereinafter referred to as the mold surface side) of the alloy flakes, but as a whole a moderately fine and homogeneous structure brought about by rapid solidification Can be obtained.
[0008]
As described above, in the RTB-based alloy cast by the SC method, since the R-rich phase is finely dispersed and the generation of α-Fe is suppressed, when producing a sintered magnet, The homogeneity of the R-rich phase in the final magnet is enhanced, and the harmful effects of α-Fe on pulverization and magnetism can be prevented. As described above, the RTB-based alloy ingot cast by the SC method has an excellent structure for producing a sintered magnet. However, as the properties of magnets have improved, it has been increasingly required to improve the homogeneity of the structure of the raw material alloy ingot.
[0009]
Therefore, for example, Japanese Unexamined Patent Publication No. 10-317110 discloses a sintered magnet having good magnet properties by reducing the area ratio of chill crystals on the mold surface side of a cast RTB-based alloy to 5% or less. Are disclosed. Since the chill crystal part becomes a fine powder having a particle size of 1 μm or less in the pulverizing step, it is considered that the particle size distribution of the alloy powder is disturbed and magnetism is deteriorated.
[0010]
Previously, the present inventors studied the relationship between the structure of the cast RTB-based alloy ingot and the behavior during hydrogen crushing or fine pulverization. As a result, the particle size of the alloy powder for the sintered magnet was determined. It has been found that it is more important to control the dispersion state of the R-rich phase than the crystal grain size of the alloy lump in order to control the uniformity (Japanese Patent Application No. 2001-393889). The volume ratio of the chill crystals in the alloy ingot is actually several percent or less, and the dispersion state of the R-rich phase generated on the mold surface side in the alloy ingot is extremely finer than the adverse effect of the chill crystals. It has been found that the region (fine R-rich phase region) has a greater effect in controlling the particle size of the magnet powder. In other words, even if the number of chill crystals in the RTB-based alloy ingot is reduced depending on the composition of the alloy ingot and the manufacturing conditions, the volume ratio of the fine R-rich phase region may exceed 50% in some cases. It was confirmed that the rich phase region disturbed the particle size distribution of the alloy powder for magnets, and that it was necessary to reduce the fine R rich phase region in order to improve the magnet properties.
[0011]
Then, the casting conditions in the SC method, especially the surface condition of the casting rotary roll, were changed, and the volume ratio at which a fine R-rich phase region was formed in the RTB-based alloy flakes was compared. Then, they found that there was a relationship between the surface roughness of the surface of the alloy flakes on the mold surface side and the volume ratio at which the fine R-rich phase region was formed, and the alloy flake having a fine R-rich phase region of 20% or less and having excellent structure homogeneity Made possible.
[0012]
[Problems to be solved by the invention]
It is possible to some extent to reduce the fine R-rich phase region and homogenize the structure by the method disclosed in Japanese Patent Application No. 2001-393889. However, since there are various factors that affect the structure other than the surface state of the casting roll, it is difficult to control all of these factors in the actual production of an RTB-based alloy. In some cases, a fine R-rich phase region is generated. Therefore, the present invention produces an alloy ingot having a structure excellent in homogeneity by suppressing the generation of a fine R-rich phase region in a cast RTB-based alloy ingot more effectively than the conventional method. Provide a method. It is another object of the present invention to further increase the homogeneity of the distribution of the R-rich phase in the magnet and to provide a rare-earth magnet having excellent magnet properties.
[0013]
[Means for Solving the Problems]
The present inventors changed the surface condition of the rotary casting roll in the SC method, and compared the volume ratio at which a fine R-rich phase region was formed in the RTB-based alloy flakes. Then, it was found that, in addition to the surface roughness of the mold surface side of the alloy flake, the shape of the irregularities and the volume ratio at which a fine R-rich phase region is formed are related. The present invention has been made by the present inventors based on the above findings.
[0014]
That is, the present invention
(12) Surface roughness of a surface having a thickness of 0.1 mm or more and 0.5 mm or less, having a plurality of linear irregularities formed so as to intersect with each other on at least one surface, and having linear irregularities. A rare earth-containing alloy flake having a ten-point average roughness (Rz) of 3 μm or more and 30 μm or less.
(2) R-T-B system used as a raw material for rare-earth magnets (where R is at least one of the rare-earth elements including Y, T is a transition metal essentially including Fe, and B is boron) The rare earth-containing alloy flake according to the above (1), which is made of an alloy.
(3) The rare earth-containing alloy flake according to (2), wherein the volume ratio of the fine R-rich phase region in the alloy is 20% or less.
(4) In the method of manufacturing a rare earth-containing alloy flake by a strip casting method (SC method), a plurality of linear irregularities are formed on a casting surface so as to intersect each other, and the surface roughness of the casting surface is a ten-point average roughness. A method for producing a rare earth-containing alloy flake, comprising using a casting roll having a height (Rz) of 3 μm or more and 30 μm or less.
(5) The rare-earth-containing alloy flake has a thickness of 0.1 mm or more and 0.5 mm or less, and has a plurality of linear irregularities formed on at least one surface so as to intersect each other. The method for producing a rare earth-containing alloy flake according to the above (4), wherein the surface having a surface has a ten-point average roughness (Rz) of 3 μm or more and 30 μm or less.
(6) A rare earth-containing alloy flake is an RTB-based alloy used as a raw material for a rare earth magnet (where R is at least one of rare earth elements including Y, T is a transition metal which essentially requires Fe, B Is boron.) The method for producing a rare earth-containing alloy flake according to the above (4) or (5), comprising an alloy.
(7) The method for producing a rare earth-containing alloy flake according to the above (6), wherein the volume ratio of the fine R-rich phase region in the alloy is 20% or less.
(8) An alloy powder for a rare-earth sintered magnet produced by subjecting the rare-earth-containing alloy flake according to the above (2) or (3) to a hydrogen crushing step and then jet milling.
(9) A rare earth sintered magnet produced by a powder metallurgy method from the alloy powder for a rare earth sintered magnet according to (8).
(10) An alloy powder for a bonded magnet manufactured by the HDDR method using the rare earth-containing alloy flake according to (2) or (3).
(11) A bonded magnet produced using the bonded magnet alloy powder according to (10).
It is.
[0015]
BEST MODE FOR CARRYING OUT THE INVENTION
FIG. 1 shows a backscattered electron image obtained by observing a cross section of a thin section of an Nd—Fe—B-based alloy (Nd 31.5% by mass) cast by a conventional SC method using a scanning electron microscope (SEM). In FIG. 1, the left side is the mold surface side, and the right side is the free surface side. The surface roughness of the alloy flakes on the mold surface side was 3.4 μm in ten-point average roughness (Rz), and the surface was almost parallel to the direction in which the linear irregularities were present.
In FIG. 1, the white portion is the Nd-rich phase (the R-rich phase is called Nd-rich phase because R is Nd), and the free surface side (opposite to the casting surface side) from the center of the alloy flake. Side surface), it forms a lamella-like pool in the thickness direction or forms a small pool with a directional shape as if the lamella were divided. However, a region in which the Nd-rich phase is extremely fine and randomly present than other portions is generated on the mold surface side, and the present inventors show that the region is a fine R-rich phase region (R Is also referred to as a fine Nd-rich phase region when the main component is Nd). This fine R-rich phase region usually starts from the mold surface side and extends toward the center. On the other hand, a portion where the fine R-rich phase region does not exist from the center to the free surface side is referred to as a normal portion here.
[0016]
In the hydrogen crushing step of the RTB-based alloy flakes during the production of the sintered magnet, hydrogen is absorbed from the R-rich phase and expands into brittle hydride. Therefore, in hydrogen crushing, fine cracks are introduced into the alloy along the R-rich phase or starting from the R-rich phase. In the subsequent pulverization step, the alloy is broken by a large number of fine cracks generated by hydrogen pulverization, so that the finer the dispersion of the R-rich phase in the alloy, the finer the particle size after the fine pulverization. Therefore, the fine R-rich phase region has a strong tendency to crack more finely than the normal part. For example, in an alloy powder manufactured from the normal part, the average particle size is about 3 μm as measured by FSSS (Fisher subsieve sizer). On the other hand, the alloy powder produced from the fine R-rich phase region contains a high proportion of fine powder of 1 μm or less, so that the particle size distribution after the fine pulverization is widened.
[0017]
JP-A-09-170055 discloses that the dispersion state of the R-rich phase in an RTB-based alloy can be controlled by controlling the cooling rate after solidification of the molten metal during casting or by heat treatment. It is described in Japanese Unexamined Patent Publication No. Hei 10-36949. However, the cooling rate after solidification or the behavior of the change of the R-rich phase inside the fine R-rich phase region due to the heat treatment is difficult to control, unlike the normal part, and the dispersion of the R-rich phase is unlikely to be coarse and remains fine. It is.
[0018]
The volume ratio of the fine R-rich phase region can be measured by the following method. FIG. 3 is a reflected electron beam image of the same field of view as that of FIG. 1, but a line is drawn at the boundary between the fine R-rich phase region and the normal part. Since the boundary between the two regions can be easily determined from the dispersion state of the R-rich phase, the area ratio of the fine R-rich phase region in the visual field can be calculated using an image analyzer. The area ratio in the cross section corresponds to the volume ratio in the alloy. In the measurement of the volume ratio of the fine R-rich phase region, the change in the amount of the fine R-rich phase region is large between flakes or within the same flake, even if the alloy flakes are simultaneously cast. Therefore, after expanding the observation field at a low magnification of about 50 to 100 times, measuring about 5 to 10 flakes and taking the average, the volume ratio of the fine R-rich phase region of the whole alloy is calculated. You can do it.
[0019]
FIG. 2 shows a reflected electron beam image of a cross section of the RTB-based alloy flake (Nd 31.5% by mass) of the present invention. In FIG. 2, the left side is the mold surface side, and the right side is the free surface side. The feature of the alloy flakes of the present invention is that, in the flakes manufactured by the strip casting method, while controlling the surface roughness on the mold surface side and forming linear irregularities on the surface so as to intersect with each other, fine flakes are obtained. This means that generation of the R-rich phase region is suppressed. The surface roughness on the mold surface side of the alloy flake shown in FIG. 2 is 3.2 μm, which is almost the same as the alloy flake of FIG. However, in the alloy flake of the present invention, the fine R-rich phase region does not exist on the mold surface side, and the dispersion state of the R-rich phase is extremely uniform from the mold surface to the free surface.
[0020]
The relationship between the surface roughness of the mold flank surface of the alloy flake manufactured by the strip casting method, the state of linear irregularities, and the fine R-rich phase region can be explained as follows.
In order for the surface of the alloy flakes to be smooth on the mold side, it is necessary that the surface of the rotating roll for casting be smooth and have good wettability with the molten alloy. In such a state, heat transfer from the molten metal to the mold is extremely good (the heat transfer coefficient is large), and the mold surface side of the alloy is excessively rapidly cooled. It is considered that the fine R-rich phase region tends to be generated when the heat transfer coefficient between the mold and the molten metal is large and the mold surface of the alloy is excessively rapidly cooled.
[0021]
On the other hand, if fine irregularities are formed on the surface of the casting rotary roll, the molten metal cannot completely enter the fine irregularities on the surface of the casting rotary roll due to the viscosity of the molten alloy. The transfer coefficient decreases. As a result, it is considered that the mold surface side of the alloy is not excessively rapidly cooled, and the generation of the fine R-rich phase region can be suppressed. Here, when the surface roughness of the surface of the rotary casting roll is increased, the surface roughness of the alloy flake is more or less transferred to the mold surface side of the alloy flake, so that the surface roughness of the mold surface of the alloy flake naturally increases. The reason why the formation of the R-rich phase is suppressed due to the alloy flake having an appropriate surface roughness on the mold side is presumed to be due to the suppression of excessive heat transfer when the molten metal solidifies as described above. Is done.
[0022]
Here, if attention is further paid to the shape of the irregularities, if the shapes are linear that do not intersect with each other, each of the contact portion and the non-contact portion between the molten metal and the roll tends to exist continuously along the linear irregularities. There is. Therefore, the internal structure also tends to have continuity along the irregularities. In this case, when a fine R-rich phase region is generated in a certain linear uneven portion for some reason, there is a high risk that a fine R-rich phase region is generated over the entire linear uneven portion.
[0023]
However, if the linear irregularities cross each other, not only the surface irregularities are divided, but also the continuity of the internal structure of the alloy is interrupted at the intersections. Further, in this intersection, the linear convex part is always cut by the concave part. It is considered that the convex portion has a good heat transfer coefficient due to good contact between the molten metal and the surface of the casting roll, so that a fine R-rich phase region is easily generated by rapid solidification. Therefore, even if a fine R-rich phase is generated, there is an effect of preventing its expansion.
[0024]
In the method for producing a rare earth-containing alloy flake by the strip casting method of the present invention, a plurality of linear irregularities are formed on the casting surface so as to intersect each other, and the surface roughness of the casting surface is the ten-point average roughness (Rz). And a casting roll having a thickness of 3 μm or more and 30 μm or less is used. As a result, at least one surface has a plurality of linear irregularities formed so as to intersect with each other, and the surface having the linear irregularities has a ten-point average roughness (Rz) of 3 μm or more and 30 μm or more. The following rare earth-containing alloy flakes can be produced. According to the present invention, even if the surface roughness is smaller than that of the conventional method, generation of a fine R-rich phase can be suppressed, and the effect of homogenizing the structure can be obtained. Further, when the surface roughness of the casting rotary roll is reduced, the amount of cutting when adjusting the roll surface can be reduced, so that the life of the casting rotary roll can be increased. Further, since it is hardly affected by the difference in the surface roughness, it is also possible to simplify the management standard of the roll surface state.
[0025]
Although the conventional SC method contained alloy flakes having a homogeneous structure as shown in FIG. 2 to some extent, flakes containing a large amount of fine R-rich phase regions as shown in FIG. 1 are also generated at the same time. As a result, there was a problem in the homogeneity of the structure throughout the alloy. Such variations in the alloy structure produced by the conventional SC method are considered to be caused by subtle differences in the contact state between the roll surface and the melt, such as the surface state of the casting roll, the supply state of the melt, and the atmosphere. Can be The irregularities formed on the surface of the rotary casting roll suppress excessive heat transfer when the molten metal solidifies, and can suppress the generation of the fine R-rich phase region with good reproducibility.
Furthermore, in the present invention, the effect of suppressing the generation of the fine R-rich phase region is enhanced by making the shape of the irregularities on the surface of the casting rotating roll into a line that intersects with each other. I got it. As a result, the yield of alloy flakes having a homogeneous structure as shown in FIG. 2 could be further increased.
[0026]
Further, details of the present invention will be described.
(1) Strip casting method (SC method)
The present invention relates to a rare earth-containing alloy flake. Here, casting of an RTB-based alloy by a strip casting method will be described.
FIG. 4 is a schematic view of an apparatus used for casting by the strip casting method. Usually, the RTB-based alloy is melted using the refractory crucible 1 in a vacuum or an inert gas atmosphere due to its active properties. After the melt of the melted alloy is held at 1350-1500 ° C. for a predetermined time, the casting rotating roll 3 whose inside is water-cooled through a tundish 2 provided with a rectifying mechanism and a slag removing mechanism as necessary. Supplied to The supply speed of the molten metal and the number of rotations of the rotating roll are appropriately controlled in accordance with the required alloy thickness. Generally, the rotation speed of the rotating roll is about 1 to 3 m / s in terms of peripheral speed. Copper or a copper alloy is suitable for the material of the casting roll because of its good thermal conductivity and easy availability. Depending on the material of the rotating roll and the surface condition of the roll, metal easily adheres to the surface of the casting rotating roll, so if a cleaning device is installed as necessary, the quality of the R-T-B alloy is stable. I do. The alloy 4 solidified on the rotating roll separates from the roll on the opposite side of the tundish and is collected in the collection container 5. By providing a heating and cooling mechanism in this collection container, the state of the R-rich phase structure in the normal part can be controlled.
[0027]
The thickness of the rare earth-containing alloy flake of the present invention is preferably 0.1 mm or more and 0.5 mm or less. If the thickness of the alloy flakes is less than 0.1 mm, the solidification rate will increase excessively, the crystal grain size will be too fine, and it will be close to the finely pulverized particle size in the magnetizing process, so the magnet orientation rate and magnetization There is a problem of inviting. On the other hand, if the thickness of the alloy flake is more than 0.5 mm, problems such as a decrease in the dispersibility of the Nd-rich phase due to a decrease in the solidification rate and precipitation of α-Fe are caused.
[0028]
(2) Surface roughness of the casting surface of the casting roll
In the present invention, when casting an RTB-based alloy by the strip casting method, the surface roughness of the casting surface of the rotary casting roll is set to 3 μm or more and 30 μm or less in ten-point average roughness (Rz).
Here, the surface roughness is measured under the conditions shown in JIS B 0601 “Definition and Display of Surface Roughness”, and the ten-point average roughness (Rz) is also defined therein. Specifically, first, a curve (roughness curve) obtained by removing a surface undulation component longer than a predetermined wavelength by a phase-compensating high-pass filter or the like from a cut (cross-sectional curve) obtained by cutting a plane perpendicular to the measurement surface. Ask. From the roughness curve, a reference length is extracted in the direction of the average line, and from the average line of the extracted portion, the average value of the absolute values of the altitudes (Yp) of the highest to fifth peaks and the lowest value The sum of the absolute values of the altitudes (Yv) of the valley bottoms from the valley bottom to the fifth is called ten-point average roughness (Rz). For measurement parameters such as the reference length, standard values for the surface roughness are specified in the JIS.
The surface roughness of the alloy flakes on the mold surface side may fluctuate greatly. At least five flakes should be measured and the average value should be used.
[0029]
(3) The shape of the irregularities on the casting surface of the rotary casting roll
In the present invention, the surface roughness of the casting surface is mainly caused by a plurality of linear irregularities formed on the surface, and the linear irregularities are formed on the surface so as to cross each other.
When the unevenness is linear, each of the contact portion and the non-contact portion between the molten metal and the roll tends to exist continuously along the linear unevenness. Therefore, the internal structure also tends to have continuity along the irregularities. In this case, if a fine R-rich phase region is generated in a certain linear uneven portion for some reason, there is a high risk that a fine R-rich phase region is generated over the entire linear portion.
However, when the linear irregularities cross each other, not only the surface irregularities are divided, but also the continuity of the internal structure is interrupted at the intersections, and even if a fine R-rich phase is generated, This has the effect of preventing its expansion.
[0030]
In the present invention, the effect of homogenizing the tissue can be obtained even if the surface roughness is relatively small, from 3 μm to 30 μm in terms of the ten-point average roughness (Rz), due to the effect of linear irregularities existing so as to cross each other. Can be obtained.
However, when the surface roughness is 3 μm or less, the effect of unevenness cannot be obtained, and the contact between the surface of the casting roll and the molten alloy is good, so that the heat transfer coefficient increases. As a result, a fine R-rich phase region is easily generated in the alloy.
[0031]
If the surface roughness of the casting rotary roll exceeds 30 μm, the solidified alloy flakes may bite into the roll surface and become difficult to peel off, which may cause troubles such as breaking the tundish. Therefore, the surface roughness of the casting roll is 30 μm or less.
[0032]
(4) Surface roughness and irregular shape of rare earth-containing alloy flakes
In the present invention, the surface roughness of at least one surface of the rare-earth-containing alloy flake is characterized by a ten-point average roughness (Rz) of 3 μm or more and 30 μm or less, and this surface roughness is mainly formed on the surface. It is provided by a plurality of linear irregularities, and the linear irregularities are formed on the surface so as to cross each other.
The surface on which the above roughness is formed is a mold surface side surface where solidification starts when casting by the strip casting method, and is a surface on which the roughness of the surface of the rotating roll is reflected. As described above, when the surface roughness of the surface is 3 μm or less, the volume ratio at which the fine R-rich phase region is generated becomes large, and the dispersion state of the R-rich phase in the alloy becomes uneven. As a result, the particle size distribution of the alloy powder after the pulverization in the manufacturing process of the sintered magnet is widened, and the characteristics of the magnet are deteriorated. On the other hand, if it is 30 μm or more, a problem is likely to occur in the process of casting the alloy. In the present invention, the surface roughness of one side of the alloy flake is 3 μm or more and 30 μm or less.
[0033]
(5) Volume fraction of fine R-rich phase region in alloy
In the present invention, the volume ratio of the fine R-rich phase region in the RTB-based alloy is 20% or less. As a result, in the manufacturing process of the sintered magnet, the particle size distribution of the alloy powder after the pulverization becomes narrow and uniform, so that it is possible to obtain a homogeneous sintered magnet having no variation in characteristics.
[0034]
(6) Alloy powder for rare earth sintered magnet and method for producing rare earth sintered magnet
From the rare-earth-containing alloy flakes composed of the R-T-B-based alloy as the raw material for magnets cast according to the present invention, high-performance anisotropic sintered magnets can be manufactured through the steps of pulverization, molding and sintering. it can.
[0035]
The pulverization of the alloy flake is usually performed in the order of hydrogen pulverization and fine pulverization, and an alloy powder of about 3 μm (FSSS) is produced. Here, the hydrogen disintegration is divided into a pre-process hydrogen storage process and a post-process dehydrogenation process. In the hydrogen storage step, hydrogen is stored mainly in the R-rich phase of the alloy flakes in a hydrogen gas atmosphere at a pressure of 266 hPa to 0.3 MPa, and the R-rich phase is expanded by the R-hydride generated at this time. By taking advantage of this, the alloy flake itself is finely divided or innumerable fine cracks are generated. This hydrogen storage is carried out at a temperature in the range of room temperature to about 600 ° C. In order to increase the volume expansion of the R-rich phase and efficiently split it, the pressure of the hydrogen gas atmosphere should be increased and the room temperature to about 100 ° C. It is preferable to carry out within the range. The preferred processing time is one hour or more. Since the R-hydride produced in this hydrogen storage step is unstable in the atmosphere and easily oxidized, after the hydrogen storage treatment, dehydration is performed by holding the alloy flake in a vacuum at about 200 to 600 ° C. and 1.33 hPa or less. It is preferable to perform elementary treatment. By this treatment, it is possible to change to R-hydride which is stable in the atmosphere. The preferred treatment time of the dehydrogenation treatment is 30 minutes or more. When atmosphere control for preventing oxidation is performed in each step from hydrogen storage to sintering, dehydrogenation treatment can be omitted.
[0036]
The RTB-based alloy flakes produced by the strip casting method of the present invention are characterized in that the R-rich phase is uniformly dispersed. The preferred average value of the interval between the R-rich phases depends on the particle size in the manufacturing process of the magnet, but is generally 3 μm to 8 μm. In hydrogen crushing, cracks are introduced into the alloy along the R-rich phase or starting from the R-rich phase. Therefore, the effect of the R-rich phase uniformly and finely dispersed in the alloy can be maximized by crushing hydrogen and then finely pulverizing, and the alloy powder having a very narrow particle size distribution can be efficiently produced. It is possible to produce. When a sintered magnet is manufactured without performing the hydrogen crushing step, the characteristics of the manufactured sintered magnet are inferior. (M. Sagawa et al. Proceeding of the 5th International Conference on Advanced Materials, Beijing China (1999))
[0037]
Fine pulverization refers to pulverizing an RTB-based alloy flake to about 3 μm (FSSS). As a pulverizing device for fine pulverization, a jet mill device is most suitable because of high productivity and a narrow particle size distribution. By using the alloy flake having a small fine R-rich phase region of the present invention, an alloy powder having a narrow particle size distribution can be produced with high efficiency and high stability.
The atmosphere during the pulverization is an inert gas atmosphere such as an argon gas or a nitrogen gas. 2% by mass or less, preferably 1% by mass or less of oxygen may be mixed into these inert gases. Thereby, the pulverization efficiency is improved, and the oxygen concentration of the pulverized alloy powder can be set to 1000 to 10000 ppm, and the alloy powder can be appropriately stabilized. At the same time, abnormal growth of crystal grains when sintering the magnet can be suppressed.
[0038]
When molding the above alloy powder in a magnetic field, to reduce the friction between the alloy powder and the inner wall of the mold, and also to reduce the friction between the powders and improve the orientation, the alloy powder such as zinc stearate is used. It is preferable to add a lubricant. A preferable addition amount is 0.01 to 1% by mass. The lubricant may be added before or after the pulverization, but it is preferable that the lubricant is sufficiently mixed using an V-type blender or the like in an inert gas atmosphere such as an argon gas or a nitrogen gas before molding in a magnetic field.
[0039]
The RTB-based alloy powder pulverized to about 3 μm (FSSS) is press-molded by a molding machine in a magnetic field. The mold is made by combining a magnetic material and a non-magnetic material in consideration of the direction of the magnetic field in the cavity. Molding pressure is 0.5-2t / cm 2 Is preferred. The magnetic field in the cavity during molding is preferably 5 to 20 kOe. The atmosphere at the time of molding is preferably an inert gas atmosphere such as an argon gas or a nitrogen gas. However, in the case of the above-mentioned oxidation-resistant powder, it can be performed in the air.
Molding can also be performed by cold isostatic pressing (CIP) or pseudo isostatic pressing (RIP: Rubber Isostatic Press) using a rubber mold. CIP and RIP are hydrostatically compressed, so that the orientation is not disturbed at the time of molding, the orientation ratio can be increased as compared with the die molding, and the maximum magnetic energy product can be increased.
[0040]
Sintering of the molded body is performed at 1000 to 1100 ° C. The sintering atmosphere is an argon gas atmosphere or 1.33 × 10 -2 A vacuum atmosphere of hPa or less is preferable. The holding time at the sintering temperature is preferably 1 hour or more. During sintering, it is necessary to remove as much as possible the lubricant contained in the molded body and the hydrogen contained in the alloy powder before reaching the sintering temperature. Preferred conditions for removing the lubricant are 1.33 × 10 -2 To hold at 300 to 500 ° C. for 30 minutes or more in a vacuum of hPa or less or in a reduced pressure Ar flow atmosphere. The preferable conditions for removing hydrogen are 1.33 × 10 -2 It is to hold at 700 to 900 ° C. for 30 minutes or more in a vacuum of hPa or less.
[0041]
After the sintering is completed, heat treatment can be performed at 500 to 650 ° C. as needed to improve the coercive force of the sintered magnet. A preferable atmosphere in this case is an argon gas atmosphere or a vacuum atmosphere, and a preferable holding time is 30 minutes or more.
[0042]
In addition, the RTB-based alloy flakes for rare earth magnets, in which the generation of the fine R-rich region produced in the present invention is suppressed, can be suitably used for producing bonded magnets in addition to sintered magnets. Hereinafter, a case where a bonded magnet is produced from the alloy flake for a rare earth magnet of the present invention will be described.
[0043]
The RTB-based alloy flake of the present invention is first heat-treated as necessary. The purpose of the heat treatment is to remove α-Fe from the alloy and to coarsen the crystal grains. To prepare an alloy powder for a bonded magnet, an HDDR (Hydrogenation Dissorption Decomposition Recombination) process is performed. However, α-Fe present in the alloy cannot be erased in the HDDR process, causing a decrease in magnetism. Become. Therefore, α-Fe needs to be erased before performing the HDDR process.
[0044]
The average particle size of the alloy powder for the bonded magnet is 50 to 300 μm, which is much larger than that of the alloy powder for the sintered magnet. In the HDDR method, the crystal orientation of the original alloy coincides with the orientation of the recombined submicron crystal grains with a certain distribution. Therefore, if two or more crystal grains having different crystal orientations in the alloy flakes of the raw material are included in one bonded magnet alloy powder, the alloy powder will include a region in which the crystal orientation is greatly different, The orientation ratio of the magnet decreases, and the maximum magnetic energy product decreases. To avoid this, the crystal grain size in the alloy flakes is preferably larger. In an alloy cast by a rapid solidification method such as a strip casting method, the crystal grain size tends to be relatively small. Therefore, the coarsening of the crystal grains by heat treatment is effective in improving the properties of the bonded magnet.
[0045]
There are many reports on a method for producing an alloy powder for a bonded magnet by the HDDR method (for example, T. Takeshita et al, Proc. 10th Int. Works on RE magnets and their application, Kyoto, Vol. 1989, Vol. 1891). ). The production of the alloy powder by the HDDR method is performed as follows.
[0046]
When the raw RTB alloy flakes are heated in a hydrogen atmosphere, the magnetic phase R-T 2 T 14 B phase is α-Fe, RH 2 , Fe 2 Decomposes into three phases B. Then, at the same temperature, when the atmosphere is switched to an inert gas atmosphere or a vacuum atmosphere to remove hydrogen, the decomposed phase has a crystal grain size of about submicron. 2 T 14 Recombines into phase B. At this time, if the composition of the alloy and the processing conditions are appropriately controlled, each recombined R 2 T 14 B-phase easy axis (R 2 T 14 B-phase C axis) is the value of R in the raw material alloy before decomposition. 2 T 14 The anisotropic magnet powder becomes almost parallel to the C-axis of the B phase, and the direction of the axis of easy magnetization of each fine crystal grain is aligned.
[0047]
The alloy that has been subjected to the HDDR treatment may be pulverized to about 50 to 300 μm to obtain an alloy powder, and then mixed with a resin and subjected to compression molding, injection molding, or the like to form a bonded magnet.
[0048]
The fine R-rich phase region has a strong tendency to be pulverized during the HDDR treatment as in the above-described hydrogen crushing treatment. The properties of the magnetic powder according to the HDDR method decrease as the particle size decreases. Therefore, the RTB-based alloy of the present invention in which the generation of the fine R-rich phase is suppressed can be suitably used for the production of the magnetic powder for the bonded magnet in the HDDR process.
[0049]
Recently, in the production of rare earth alloys by the SC method, it has been reported that the homogeneity of the alloy structure was improved by controlling the surface roughness Sm / Ra and Sm of the outer peripheral surface of the casting roll to a certain range (particularly). Japanese Unexamined Patent Application Publication No. 2002-59245 and Japanese Unexamined Patent Application Publication No. 9-1296. However, they are intended to suppress a change in the structure in the width direction of the strip and to prevent a decrease in cooling rate at the end of the strip. Further, the shape of the unevenness that causes the surface roughness is not particularly defined.
On the other hand, the present invention is intended to suppress the structural change in the thickness direction of the alloy flakes, that is, from the roll surface side to the free surface side, and to homogenize the structure, and introduce a fine R-rich phase region as a measure of homogeneity. However, the specific volume ratio is clearly shown, and is completely different from the invention disclosed in JP-A-2002-59245 or JP-A-9-1296.
[0050]
【Example】
(Example 1)
The alloy composition is such that Nd: 31.5 mass%, B: 1.00 mass%, Co: 1.0 mass%, Al: 0.30 mass%, Cu: 0.10 mass%, and the balance iron. Using a alumina crucible, a raw material containing a mixture of neodymium, metal, ferroboron, cobalt, aluminum, copper, and iron is melted in a high-frequency melting furnace in an atmosphere of 1 atm with argon gas, and the molten metal is strip cast. An alloy flake was produced by casting.
The diameter of the casting roll was 300 mm, the material was pure copper, the inside was water-cooled, and the surface roughness of the casting surface was adjusted to 4.0 μm as a ten-point average roughness (Rz). The surface roughness at this time was mainly caused by irregularities in a random direction such as intersecting linear lines provided on the casting surface. The peripheral speed of the roll during casting was 1.0 m / s, and alloy flakes having an average thickness of 0.30 mm were produced.
[0051]
The surface roughness of the mold flake side of the obtained alloy flake was 4.6 μm in ten-point average roughness (Rz), and had linear irregularities crossing each other. After embedding and polishing 10 alloy flakes, a reflected electron beam image (BEI) of each alloy flake was taken at a magnification of 100 times with a scanning electron microscope (SEM). When the photographed image was taken into an image analyzer and measured, the volume ratio of the fine R-rich phase region was 3% or less.
[0052]
(Comparative Example 1)
The raw materials were blended in the same composition as in Example 1, and melting and casting by the SC method were performed in the same manner as in Example 1. However, the surface roughness of the casting roll is 4.0 μm in ten-point average roughness (Rz), and has linear irregularities substantially parallel to the roll rotation direction. Substantially absent.
As a result of evaluating the obtained alloy flakes in the same manner as in Example 1, the surface roughness of the mold surface side was 4.5 μm in ten-point average roughness (Rz), and the volume ratio of the fine R-rich phase region was: 25%.
[0053]
(Comparative Example 2)
The raw materials were blended in the same composition as in Example 1, and melting and casting by the SC method were performed in the same manner as in Example 1. However, the surface roughness of the surface of the rotary casting roll was 100 μm in terms of a ten-point average roughness (Rz), and the surface roughness was mainly caused by linear irregularities crossing each other as in Example 1. It had been.
Then, during the casting, a part of the metal made one round without separating from the roll, came into contact with the tundish, and its front end was broken, so the casting operation was stopped.
[0054]
Next, an example in which a sintered magnet is manufactured will be described.
(Example 2)
The alloy flake obtained in Example 1 was pulverized with hydrogen and pulverized with a jet mill. The condition of the hydrogen storage step, which is a step before the hydrogen disintegration step, was kept at 100% hydrogen atmosphere and 2 atm for 1 hour. The temperature of the metal piece at the start of the hydrogen storage reaction was 25 ° C. The condition of the dehydrogenation step, which is a post-step, was maintained at 500 ° C. for 1 hour in a vacuum of 0.133 hPa. To this powder, 0.07% by mass of zinc stearate powder was added, mixed sufficiently in a 100% nitrogen atmosphere by a V-type blender, and then finely pulverized by a jet mill. The pulverizing atmosphere was a nitrogen atmosphere mixed with 4000 ppm of oxygen. Thereafter, the mixture was again sufficiently mixed in a 100% nitrogen atmosphere by a V-type blender. The obtained powder had an oxygen concentration of 2500 ppm, and from the analysis of the carbon concentration of the powder, it was calculated that the zinc stearate powder mixed with the powder was 0.05% by mass. Moreover, as a result of measuring with a laser diffraction type particle size distribution analyzer, the average particle size D50 was 5.00 μm, D10 was 1.98 μm, and D90 was 8.51 μm.
[0055]
Next, the obtained powder was press-molded in a 100% nitrogen atmosphere using a molding machine in a horizontal magnetic field. Molding pressure is 1.2t / cm 2 And the magnetic field in the cavity of the mold was 15 kOe. 1.33 × 10 -5 Hold at 500 ° C. for 1 hour in a vacuum of hPa, then 1.33 × 10 -5 After maintaining at 800 ° C. for 2 hours in a vacuum of hPa, the pressure was further increased to 1.33 × 10 3 -5 It was sintered at 1050 ° C. for 2 hours in a vacuum of hPa. The sintered density is 7.5 g / cm 3 This is a sufficient density. Further, this sintered body was heat-treated at 560 ° C. for 1 hour in an argon atmosphere to produce a sintered magnet.
[0056]
Table 1 shows the results of measuring the magnetic properties of the sintered magnet using a DC BH curve tracer. Table 1 also shows the oxygen concentration and particle size of the fine powder of the raw material of the sintered magnet.
[0057]
(Comparative Example 3)
The alloy flake obtained in Comparative Example 1 was pulverized in the same manner as in Example 2 to obtain a fine powder. Further, through the same molding and sintering steps as in Example 2, a sintered magnet was produced.
[0058]
Table 1 shows the results of measuring the magnetic properties of the sintered magnet produced in Comparative Example 3 using a direct current BH curve tracer. Table 1 also shows the oxygen concentration and particle size of the fine powder of the raw material for the sintered magnet of Comparative Example 3.
[0059]
[Table 1]
Figure 2004043921
[0060]
As shown in Table 1, since D10 is smaller in Comparative Example 3 than in Example 2, it can be seen that the ratio of very fine powder smaller than about 1 μm is large. Such very fine particles are easily oxidized, and the oxygen concentration of the fine powder in Comparative Example 3 is slightly higher than that in Example 2. It is considered that the magnetic properties of the magnet of Comparative Example 3 are lower than those of Example 2 mainly due to the increase in oxygen concentration and the heterogeneity of the crystal structure.
[0061]
Next, an example in which a bonded magnet is manufactured will be described.
(Example 3)
The raw materials were blended such that the alloy composition was 28.5% by mass of Nd, 1.00% by mass of B, 10.0% by mass of Co, 0.5% by mass of Ga, and the balance of iron. Under the conditions, an alloy flake was cast by the SC method.
As a result of evaluating the obtained alloy flakes in the same manner as in Example 1, the surface roughness of the mold surface side was 4.3 μm in ten-point average roughness (Rz), and the volume ratio of the fine R-rich phase region was 3% or less. And α-Fe was not contained.
[0062]
After holding the above alloy flakes in hydrogen at 1 atm at 820 ° C. for 1 hour, HDDR treatment was performed in which the alloy flakes were held at the same temperature in vacuum for 1 hour. The obtained alloy powder was pulverized with a brown mill to 150 μm or less, added with 2.5% by mass of an epoxy resin, and compression-molded by applying a magnetic field of 1.5 T to obtain a bonded magnet. Table 1 shows the magnetic properties of the obtained bonded magnet.
[0063]
(Comparative Example 4)
The raw materials were blended in the same composition as in Example 3, and melting and casting by the SC method were performed in the same manner as in Comparative Example 1. The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface side was 4.8 μm in ten-point average roughness (Rz), and the volume ratio of the fine R-rich phase region was 30%. Met.
[0064]
Next, using the alloy flake obtained in Comparative Example 4, a bonded magnet was produced in the same manner as in Example 3. Table 1 shows the magnetic properties of the obtained bonded magnet.
[0065]
Table 1 shows that the bonded magnets of Example 3 and Comparative Example 4 have excellent magnetic properties of Example 3. In Comparative Example 4, since the volume ratio of the fine R-rich region is high and the amount of relatively fine particles of 50 μm or less after HDDR treatment or pulverization is large, it can be estimated that the magnetism is low.
[0066]
【The invention's effect】
The RTB-based alloy flake of the present invention has a small volume ratio of the fine R-rich region, and the homogeneity of the dispersed state of the R-rich phase in the alloy is even better than the alloy flake manufactured by the conventional SC method. It is. Therefore, the sintered magnet manufactured from the RTB-based alloy flake of the present invention and the bonded magnet manufactured by the HDDR method exhibit superior magnet properties as compared with the conventional magnet.
[Brief description of the drawings]
FIG. 1 is a view showing a cross-sectional structure of a rare-earth magnet alloy flake containing a fine R-rich phase manufactured by a conventional SC method.
FIG. 2 is a view showing a cross-sectional structure of a rare-earth magnet alloy flake according to the present invention.
FIG. 3 is a diagram in which a line is drawn at a boundary between a fine R-rich region and a normal portion in the sectional structure of FIG. 1;
FIG. 4 is a schematic view of a casting apparatus using a strip casting method.
[Explanation of symbols]
1 refractory crucible
2 Tundish
3 Rotating roll for casting
4 Alloy
5 Collection container

Claims (11)

厚さが0.1mm以上0.5mm以下であり、少なくとも片側の表面に互いに交差するように形成された複数の線状の凹凸を有し、線状の凹凸を有する面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下であることを特徴とする希土類含有合金薄片。It has a thickness of 0.1 mm or more and 0.5 mm or less, has a plurality of linear irregularities formed so as to intersect each other on at least one surface, and the surface having the linear irregularities has a sufficient surface roughness. A rare earth-containing alloy flake having a point average roughness (Rz) of 3 μm or more and 30 μm or less. 希土類磁石用原料として使用されるR−T−B系(但し、RはYを含む希土類元素のうち少なくとも1種、TはFeを必須とする遷移金属、Bは硼素である。)合金からなることを特徴とする請求項1に記載の希土類含有合金薄片。An R-T-B-based alloy used as a raw material for a rare earth magnet (where R is at least one of rare earth elements including Y, T is a transition metal essentially containing Fe, and B is boron). The rare earth-containing alloy flake according to claim 1, wherein: 合金中の微細Rリッチ相領域の体積率が20%以下であることを特徴とする請求項2に記載の希土類含有合金薄片。The rare earth-containing alloy flake according to claim 2, wherein the volume ratio of the fine R-rich phase region in the alloy is 20% or less. ストリップキャスト法(SC法)による希土類含有合金薄片の製造方法において、鋳造面に互いに交差するように複数の線状の凹凸が形成され、該鋳造面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下である鋳造用回転ロールを用いることを特徴とする希土類含有合金薄片の製造方法。In a method of manufacturing a rare earth-containing alloy flake by a strip casting method (SC method), a plurality of linear irregularities are formed so as to intersect with each other on a casting surface, and the surface roughness of the casting surface is a ten-point average roughness (Rz). A) using a casting roll having a size of 3 μm or more and 30 μm or less. 希土類含有合金薄片が、厚さが0.1mm以上0.5mm以下であり、少なくとも片側の表面に互いに交差するように形成された複数の線状の凹凸を有し、線状の凹凸を有する面の表面粗さが十点平均粗さ(Rz)で3μm以上30μm以下であることを特徴とする請求項4に記載の希土類含有合金薄片の製造方法。The rare earth-containing alloy flake has a thickness of 0.1 mm or more and 0.5 mm or less, and has a plurality of linear irregularities formed so as to intersect each other on at least one surface, and a surface having linear irregularities. 5. The method for producing a rare earth-containing alloy flake according to claim 4, wherein the surface roughness is 10 μm to 30 μm in terms of ten-point average roughness (Rz). 6. 希土類含有合金薄片が、希土類磁石用原料として使用されるR−T−B系(但し、RはYを含む希土類元素のうち少なくとも1種、TはFeを必須とする遷移金属、Bは硼素である。)合金からなることを特徴とする請求項4または5に記載の希土類含有合金薄片の製造方法。Rare earth-containing alloy flakes are used as a rare earth magnet raw material in an RTB system (where R is at least one of the rare earth elements including Y, T is a transition metal in which Fe is essential, and B is boron. 6. The method for producing a rare earth-containing alloy flake according to claim 4, wherein the alloy flake is made of an alloy. 合金中の微細Rリッチ相領域の体積率が20%以下であることを特徴とする請求項6に記載の希土類含有合金薄片の製造方法。The method for producing a rare earth-containing alloy flake according to claim 6, wherein the volume ratio of the fine R-rich phase region in the alloy is 20% or less. 請求項2又は3に記載の希土類含有合金薄片に水素解砕工程を施し、その後にジェットミル粉砕することで作製される希土類焼結磁石用合金粉末。An alloy powder for a rare-earth sintered magnet produced by subjecting the rare-earth-containing alloy flake according to claim 2 or 3 to a hydrogen disintegration step, followed by jet milling. 請求項8に記載の希土類焼結磁石用合金粉末から粉末冶金法で製造される希土類焼結磁石。A rare earth sintered magnet produced from the alloy powder for a rare earth sintered magnet according to claim 8 by a powder metallurgy method. 請求項2又は3に記載の希土類含有合金薄片を用いて、HDDR法で製造したボンド磁石用合金粉末。An alloy powder for a bonded magnet produced by the HDDR method using the rare earth-containing alloy flake according to claim 2 or 3. 請求項10に記載のボンド磁石用合金粉末を用いて作製されるボンド磁石。A bonded magnet produced using the alloy powder for a bonded magnet according to claim 10.
JP2002205214A 2001-12-18 2002-07-15 Rare-earth-containing alloy flake, its manufacturing process, rare-earth sintered magnet, alloy powder for this, bond magnet and alloy powder for this Pending JP2004043921A (en)

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AU2002358316A AU2002358316A1 (en) 2001-12-18 2002-12-18 Alloy flake for rare earth magnet, production method thereof, alloy powder for rare earth sintered magnet, rare earth sintered magnet, alloy powder for bonded magnet and bonded magnet
US10/498,932 US7442262B2 (en) 2001-12-18 2002-12-18 Alloy flake for rare earth magnet, production method thereof, alloy powder for rare earth sintered magnet, rare earth sintered magnet, alloy powder for bonded magnet and bonded magnet
PCT/JP2002/013231 WO2003052778A1 (en) 2001-12-18 2002-12-18 Alloy flake for rare earth magnet, production method thereof, alloy powder for rare earth sintered magnet, rare earth sintered magnet, alloy powder for bonded magnet and bonded magnet
CNB028050975A CN1306527C (en) 2001-12-18 2002-12-18 Rare earth magnetic alloy sheet, its manufacturing method, sintered rare earth magnetic alloy powder, sintered rare earth magnet, metal powder for bonded magnet, and bonded magnet
US11/826,114 US7571757B2 (en) 2001-12-18 2007-07-12 Alloy flake for rare earth magnet, production method thereof, alloy powder for rare earth sintered magnet, rare earth sintered magnet, alloy powder for bonded magnet and bonded magnet

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JP2007119882A (en) * 2005-10-31 2007-05-17 Showa Denko Kk R-t-b based alloy, method for producing r-t-b based alloy sheet, fine powder for r-t-b based rare earth permanent magnet and r-t-b based rare earth permanent magnet
WO2009075351A1 (en) 2007-12-13 2009-06-18 Showa Denko K.K. R-t-b alloy, process for production of r-t-b alloy, fine powder for r-t-b rare earth permanent magnets, and r-t-b rare earth permanent magnets
US7722726B2 (en) 2004-03-31 2010-05-25 Santoku Corporation Process for producing alloy slab for rare-earth sintered magnet, alloy slab for rare-earth sintered magnet and rare-earth sintered magnet
EP2740551A4 (en) * 2011-08-03 2015-11-11 Santoku Corp Alloy flakes as starting material for rare earth sintered magnet and method for producing same

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US7722726B2 (en) 2004-03-31 2010-05-25 Santoku Corporation Process for producing alloy slab for rare-earth sintered magnet, alloy slab for rare-earth sintered magnet and rare-earth sintered magnet
US8105446B2 (en) 2004-03-31 2012-01-31 Santoku Corporation Process for producing alloy slab for rare-earth sintered magnet, alloy slab for rare-earth sintered magnet and rare-earth sintered magnet
JP2007119882A (en) * 2005-10-31 2007-05-17 Showa Denko Kk R-t-b based alloy, method for producing r-t-b based alloy sheet, fine powder for r-t-b based rare earth permanent magnet and r-t-b based rare earth permanent magnet
WO2009075351A1 (en) 2007-12-13 2009-06-18 Showa Denko K.K. R-t-b alloy, process for production of r-t-b alloy, fine powder for r-t-b rare earth permanent magnets, and r-t-b rare earth permanent magnets
EP2740551A4 (en) * 2011-08-03 2015-11-11 Santoku Corp Alloy flakes as starting material for rare earth sintered magnet and method for producing same
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