JP3721831B2

JP3721831B2 - Rare earth magnet alloy and method for producing the same

Info

Publication number: JP3721831B2
Application number: JP06463799A
Authority: JP
Inventors: 洋一広瀬; 史郎佐々木; 寛長谷川
Original assignee: Showa Denko KK
Current assignee: Resonac Holdings Corp
Priority date: 1999-03-11
Filing date: 1999-03-11
Publication date: 2005-11-30
Anticipated expiration: 2020-11-30
Also published as: JPH11315357A

Description

【０００１】
【産業上の利用分野】
本発明は希土類元素を含む磁石の原料となる原料合金およびこの原料合金の製造方法に関する。
【０００２】
【従来の技術】
最近、希土類合金系の優れた磁気特性を活かした希土類焼結磁石あるいは希土類ボンド磁石が注目されてきており、特にＲ−Ｆｅ−Ｂ系磁石において、磁気特性をさらに向上させた磁石の開発が行われている。Ｒ−Ｆｅ−Ｂ系磁石では磁性を担う強磁性相Ｒ₂ Ｆｅ₁₄Ｂ相の他に、非磁性でＮｄ等の希土類元素の濃度の高い相（Ｒリッチ相と呼ぶ）が存在し、次の様な重要な役割を果たしている。
▲１▼融点が低く、磁石化工程の焼結時に液相となり、磁石の高密度化、したがって磁化の向上に寄与する。
▲２▼粒界の凹凸をなくし、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▼Ｒリッチ相は非磁性であり主相を磁気的に絶縁することから、保磁力を高める。
したがって、Ｒリッチ相の体積率が低いか、分散状態が悪いためにＲリッチ相に覆われていない界面が存在すれば、その部分では局所的な保磁力低下によって角型性が悪化すると共に、焼結不良によって磁化も低下するため最大磁気エネルギー積の低下をもたらすことが知られている。
【０００３】
ところが、高特性磁石になるほど強磁性相であるＲ₂ Ｆｅ₁₄Ｂ相の体積率を高める必要があるため、必然的にＲリッチ相の体積率が減少し、部分的なＲリッチ相不足を生じ、十分な特性が得られない場合が多い。そこで高特性材のＲリッチ相不足による特性低下防止に関した多くの研究が報告されており、それらは大きく２つのグル−プに分けられる。
【０００４】
一つは主相Ｒ₂ Ｆｅ₁₄Ｂ相とＲリッチ相を別々の合金から供給するものであり、一般に２合金法と呼ばれている。殆どＲ₂ Ｆｅ₁₄Ｂ相単相からなる合金とＲリッチ相を生成する合金の２種類をそれぞれ微粉砕した後、適当な比率で混合、成形焼結する方法で、Ｒリッチ相の分散性を改善するためにＲリッチ相を生成する合金に多くの工夫が見られる。例えば、焼結温度での液相組成の非晶質合金を使用すれば、Ｒリッチ相よりもＦe の含有量が多いため、同じ組成の磁石を作製するのにＲリッチ相を混合するよりも主相を生成する合金との混合比率を高くでき、結果として焼結時に生成するＲリッチ相の分散性が良好となり、磁気特性向上に成功している（E.Otsuki,T.Otsuka and T.Imai,11th Internatinal Workshop on Rare Earth magnets and their Applications,vol.1,p328(1990) ）。
【０００５】
もう一つはストリップキャスティング法により、従来の金型鋳造法よりも高い冷却速度で凝固することで組織を微細化し、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ものである。合金内のＲリッチ相が微細に分散しているため、粉砕、焼結後のＲリッチ相の分散性も良好となり、磁気特性向上に成功している（特開平5-222488、特開平5-295490）。
【０００６】
【発明が解決しようとする課題】
以上のように２合金法とストリップキャスティング法によってＲリッチ相の良好な分散がもたらされ、磁気特性の向上がなされたが、これらの方法は以下のような問題を抱えている。
【０００７】
まず、前者の２合金法では、極めて高い特性が報告されているが、Ｒリッチ相を生成する非晶質合金の作製に液体超急冷を使用しなければならないが、液体超急冷は生産性が悪くコストアップをもたらす。液体超急冷を利用せず、Ｒリッチ相生成合金にＣo 系金属間化合物を使用して45MGOe以上の特性を発現した報告も存在する（楠、美濃輪、本島、電気学会論文誌A、113 巻12号、P849-853、1993）。しかし、Ｒ₂ Ｆｅ₁₄Ｂ相はα−Ｆｅと液相から包晶反応によって生成するため、ほぼＲ₂ Ｆｅ₁₄Ｂ化学量論組成での鋳造を要する主相を生成する合金では、α−Ｆｅの残存量が極めて多い。そして、α−Ｆｅは粉砕性を著しく害し、粉砕時の組成変動の原因となり、磁気特性の低下、バラツキの増加を引き起こす。そのため、主相を生成する合金は、初晶であるα−Ｆｅを消滅し、ほぼＲ₂ Ｆｅ₁₄Ｂ相単相とする必要から長時間の均質化処理を必要とする。さらに２合金法では２種類の合金粉を別々に粉砕し、均一に混合しなければならないため、従来の as cast の１合金を粉砕する方法と比較すると生産性が格段に悪く、コストの増加が避けられないため、45MGOe程度以上の付加価値が高い極めて高特性な材料にしか使用できない。
【０００８】
一方、ストリップキャスティング法も鋳造に時間がかかり、高活性の希土類元素を含む溶湯を長時間保持し、少量ずつ供給するため、ルツボ、保持炉あるいはタンディッシュと溶湯の反応により、成分が変動しやすい。また、温度を一定に保ち、定常状態で安定した鋳造を持続させるのが極めて難しく、収率が低いといった問題がある。さらに、特殊で高価な鋳造設備を必要とする等の問題もあり、コストアップは避けられない。
【０００９】
【課題を解決するための手段】
本発明者は特殊で高価な鋳造設備を使用せずに、従来の金型鋳造法によってＲリッチ相の不足、偏在による特性劣化が生じにくい合金を作製する方法について検討した結果、従来単に粒界相又はＲリッチ相と呼ばれていた、鋳造時に最後まで液相として存在するＲに富んだ部分が凝固した領域（共晶領域と呼ぶ）に注目した。その結果、その共晶部分が凝固する600 〜800 ℃辺りでの冷却速度によって組織形態が変化し、さらに磁石の保磁力、角型性とも関連する事実を見いだした。さらに、主相の結晶粒径の大きさ、共晶領域の分散状態の制御によってさらに効果が高まる事実も見いだした。本発明はこれらの知見に基づいてなされたものである。
【００１０】
すなわち本発明はＲ（Ｙを含む希土類元素のうち少なくとも１種）、Ｔ（Ｆe を必須とする遷移金属）及びＢ（硼素）を基本成分とする永久磁石の原料合金に於て、従来単に粒界相又はＲリッチ相と呼ばれていた、鋳造時に最後まで液相として存在するＲに富んだ部分が凝固した領域、すなわち共晶領域の組織を制御することにより、実質的に利用できるＲリッチ相の体積率を増加すること、さらに主相の結晶粒径、共晶領域の分散状態の制御により上記課題を解決したものである。
【００１１】
次に本発明の構成を以下に詳細に記す。
図１、図２に本発明と従来の合金の代表的な顕微鏡組織の模式図を示す。図２の従来の合金では主相１が柱状に晶出し、その周囲を最後に凝固する共晶領域２がとり囲んだ組織を呈する。また図１に示す本発明の合金ではほとんどの主相１の結晶粒の内部にも共晶領域２が存在する。図２の方がやや拡大してあり、主相１の大きさは図１の本発明合金の方が大きい。
（１）共晶領域の組織
共晶領域中のＲの含有量がより少ない相の短軸方向の大きさを３μｍ以下にしたことを特徴とする。
最後まで液相として存在するＲ（希土類）に富んだ部分は、６００〜８００℃の間に最終的には共晶反応によって幾つかの固相に変態、凝固する。例えばＮｄ−Ｆｅ−Ｂ三元系では６６５℃で共晶反応によって、液相がＮｄメタル相、Ｎｄ₂ Ｆｅ₁₄Ｂ相、ＮｄＦｅ₄ Ｂ₄ 相の３相に変態、凝固することが知られている。他の成分元素を含む場合には反応温度、生成相が変化すると思われる。本来、Ｎｄリッチ相とはＮｄ−Ｆｅ−Ｂ三元系ではＮｄメタル相を指すものであるが、共晶反応で生成したＮｄ₂ Ｆｅ₁₄Ｂ相、ＮｄＦｅ₄ Ｂ₄ 相も合わせた総称として扱っている場合も存在し、特に原料合金の組織を議論する際に多く見られる。他の成分元素を含む場合も同様である。そこで本発明では従来広くＲリッチ相、又は粒界相とも呼ばれていた、鋳造時に最後まで液相として存在するＲに富んだ部分が凝固した領域を共晶領域２と呼び、Ｒの一次固溶体であるＲリッチ相と区別して表現する。
【００１２】
図３に本発明による合金の共晶領域内の組織を拡大した模式図を示す。共晶領域２の内部には周囲よりもＲの含有量がより少ない相４が細い棒状に析出した組織を呈している。共晶領域２の内部のＲの含有量がより少ない相４とは、先に示したＲ₂ Ｆｅ₁₄Ｂ相、ＲＦｅ₄ Ｂ₄ 相であり、他の成分元素を含む場合には異なった相が存在する場合もあるが、周囲のＲ含有量の多い相３以外の相とも言える。このＲの含有量がより少ない相４の短軸方向の大きさが、磁石化工程で磁場成形用の粉末を得る微粉砕粒径程度、つまり３〜５μｍ以下であれば、これらの相４を含む粉砕粉は常にＲの含有量が多い相３と共存することになる。そのためＲの含有量が多い相３を含む粉末の比率が増加し、Ｒの含有量が多い相３の分散性の改善、焼結時の高密度化を促進、さらに焼結後の熱処理時の保磁力の温度依存性を改善し、優れた磁気特性の磁石の製造が可能となる。
【００１３】
一方、図４に従来法による合金の共晶領域内の組織を拡大した模式図を示す。従来合金ではＲ含有量のより少ない相４の大きさが、本発明の場合よりも太く、あるものは片状を呈している。本発明による合金に比較して、共晶領域２内部のＲの含有量がより少ない相４の短軸方向の大きさが微粉砕粒径程度より大きく、Ｒの含有量が多い相３を含まない粉末の比率が増加するため、Ｒの含有量が多い相３の分散性は低下し、磁気特性は低下する。Ｒの含有量がより少ない相４の微細化にはこれらの相を生成する共晶反応時の冷却速度を増加し、核生成頻度を高めつつ各生成相の粗大化を防止する方法が有効である。また、ある種の添加元素によって微細化が促進される可能性も存在する。
【００１４】
（２）主相の結晶粒径
主相であるＲ₂ Ｆｅ₁₄Ｂ結晶粒の長軸方向の大きさは５０μｍ以上、短軸方向の大きさは１０μｍ以上で、この主相の領域の体積率が７０％以上である。
凝固方向に沿って切断した図１に示す顕微鏡組織において、主相１の結晶粒径が磁場成形用の粉末を得る微粉砕粒径程度、つまり３〜５μｍ程度以下であると一つの粉砕粉中に方位の異なる２つ以上の主相が存在することになり、配向性が低下する。したがって主相１の結晶粒径は大きい方が都合が良く、長軸方向の大きさ（Ｌ）が５０μｍ以上、短軸方向の大きさ（Ｗ）が１０μｍ以上である結晶粒の体積率が７０％以上であることが好ましい。より好ましくは、長軸方向の大きさ（Ｌ）が１００μｍ以上、短軸方向の大きさ（Ｗ）が２０μｍ以上が良い。主相の各結晶粒は合金をエメリー紙で研磨した後、アルミナ、ダイアモンド等を使用してバフ研磨した面を偏光顕微鏡で観察することにより容易に識別可能である。偏光顕微鏡では磁気Ｋｅｒｒ効果により、入射した偏光が強磁性体表面の磁化方向に応じた偏光面の回転を生じて反射するため、各結晶粒から反射する偏光面の相違が明暗として観察される。
本発明の合金の偏光顕微鏡写真を見ると、主相Ｒ₂ Ｆｅ₁₄Ｂの結晶粒界が認められる。各結晶粒の結晶方向の相違が明暗となって明瞭に識別可能である。また、黒い筋状となって観察される共晶領域が主相結晶粒界だけでなく、主相Ｒ₂ Ｆｅ₁₄Ｂ結晶粒内にもその長軸方向にほぼ平行に存在している。なお、主相結晶粒内の細かい縞模様は磁区に対応している。
【００１５】
（３）共晶領域の分布
共晶領域２が主相Ｒ₂ Ｆｅ₁₄Ｂ１の結晶粒界の他に、主相１の結晶粒内にも存在する。特に共晶領域２同士の間隔（Ｄ）は主相Ｒ₂ Ｆｅ₁₄Ｂ結晶粒１の短軸方向の大きさ（Ｗ）の２分の１以下である領域の体積率が７０％以上である。
ここで共晶領域２同士の間隔（Ｄ）とは、結晶粒界の共晶領域同士であっても結晶粒内の共晶領域同士であっても良い。あるいは、また結晶粒界と結晶粒内の共晶領域の間隔であっても良い。つまり近接する共晶領域同士の間隔が主相の短軸方向の大きさの２分の１以下であれば良い。
(1) 項にこの合金を微粉砕した場合に共晶領域２内部の組織微細化によって、Ｒの含有量が多い相３を含む粉末の比率が増加し、Ｒの含有量が多い相３の分散性の改善により、焼結性並びに磁気特性の向上が可能となることを記した。そして、Ｒの含有量が多い相３の分散性はＲの含有量が多い相３を含む共晶領域２が、合金内で微細分散することでさらに良好となる。つまり共晶領域２と主相１との界面の拡大に応じて、微粉砕時のＲの含有量が多い相３を含む粉砕粉の量が増加するため、Ｒの含有量が多い相３の分散性向上をもたらす。したがって、共晶領域２は微細分散した方が都合が良く、具体的には共晶領域２同士の間隔（Ｄ）が主相のＲ₂ Ｆｅ₁₄Ｂ結晶粒１の短軸方向の大きさ（Ｗ）の２分の１以下である領域の体積率が７０％以上であることが好ましい。
【００１６】
本発明の合金では、共晶領域２が柱状晶組織では主相Ｒ₂ Ｆｅ₁₄Ｂ１の結晶粒界と結晶粒内にもその長軸方向に平行に伸長して存在し、等軸晶組織では球状に存在しているものがある。したがって、共晶領域２と主相１との界面を拡大し、Ｒの含有量が多い相３の分散性を高めるには、共晶領域２が薄く長く伸長する柱状晶組織の方が好ましい。一方、従来の鋳造法で得られる合金はストリップキャスティング法を含めて、結晶粒を構成する主相Ｒ₂ Ｆｅ₁₄Ｂ１を結晶粒界を構成するＲの含有量が多い相３が被覆している組織構造（特開平5-295490）であり、本発明による主相Ｒ₂ Ｆｅ₁₄Ｂ１の結晶粒内の複数のＲの含有量が多い相３を含む共晶領域２の存在に関する記述はなく、全く別の組織と言える。
【００１７】
(4) 製造方法
本発明の合金の製造方法について説明すると、一つは鋳造後、合金インゴットを８００℃〜１１５０℃に加熱し、８００℃から６００℃の温度域を５℃／秒以上の冷却速度で冷却することを特徴とする
もう一つは鋳造時に８００℃から６００℃の温度域を５℃／秒以上の冷却速度で冷却する。
以下、それぞれの工程について説明する。
【００１８】
本発明による共晶領域２内のＲの含有量のより少ない相４の短軸方向の大きさ（Ｗ）が３μｍ以下である組織は、鋳造後の熱処理によって生成する。つまり、一度凝固した合金を加熱して共晶領域２を再融解した後、共晶領域２が凝固する温度域での冷却速度を増加し、核生成頻度を高めつつ各生成相の粗大化を防止することで、共晶領域２内のＲの含有量がより少ない相４が微細化する。組成によって凝固温度は多少上下するが、具体的な加熱温度は共晶領域２が融解する温度よりも高い８００℃以上であり、主相Ｒ₂ Ｆｅ₁₄Ｂ相１が融解する１１５０℃以下とする必要がある。冷却は８００℃から６００℃の温度域を５℃／秒以上の冷却速度とすることが好ましく、より好ましい冷却速度は１０℃／秒以上である。また、冷却速度増加には特別な装置を必要としないガス急冷でも可能であるが、冷却効率を高めるために合金インゴットを細かく、具体的には５ｍｍ程度に粉砕することが好ましい。
【００１９】
次に、鋳造時の冷却速度を制御する方法について説明する。
共晶領域２を生成する温度域、つまり最後まで液相として存在するＲ（希土類）に富んだ部分が凝固する温度域での冷却速度を増加し、核生成頻度を高めつつ各生成相の粗大化を防止することで、共晶領域２内のＲの含有量がより少ない相４が微細化する。組成によって凝固温度は多少上下するが、具体的には８００℃から６００℃の温度域を５℃／秒以上の冷却速度で冷却することが好ましく、より好ましい冷却速度は１０℃／秒以上である。このような条件は、ストリップキャスティング法等の特殊な鋳造法を用いなくても、例えば、従来の鋳造法でもモールド比（鋳型比）を十分大きくすることによっても達成可能である。
【００２０】
モールド比は鋳型の熱容量を鋳造する合金の熱容量で除した値であり、簡単にはそれぞれの重量比、より簡単にはそれぞれの側板の厚さの比でも十分比較可能である。例えば、代表的な鋳型材である鉄と銅は単位体積当りの熱容量がほぼ等しいことから、側板厚さの比で議論すれば材質の相違も含めて比較することが可能である。モールド比の増加により、鋳型の温度上昇が抑制されるため、従来鋳型の温度上昇でインゴットの冷却効率が特に低下していた８００℃以下での冷却速度の増加が可能となる。モールド比の増加は側板の厚肉化かインゴットの簿肉化により可能である。さらに効率的な水冷機構を有することで溶湯から移動した熱を速やかに除去し、鋳型の温度上昇を防止する方法も有効であり、両者を組合わせることでより効率的な冷却が可能となる。その他、遠心鋳造法等によるインゴットの簿肉化も有効である。
【００２１】
本発明では８００℃以上、６００℃以下の冷却速度は特に規定しないが、８００℃以上については主相Ｒ₂ Ｆｅ₁₄Ｂ相１の結晶粒径、共晶領域２の分散性に影響することは良く知られている。一般的なＲ−Ｆｅ−Ｂ系磁石組成では液相から初晶α−Ｆｅが生成して、これと液相から包晶反応によってＲ₂ Ｆｅ₁₄Ｂ相１を生成するか、或は凝固速度が早い場合には包晶反応温度以下まで過冷却されて液相からＲ₂ Ｆｅ₁₄Ｂ相１を直接生成する。したがって、この包晶反応が終了するまでの冷却速度のマクロ的な合金組織への影響は特に大きく、冷却速度が大きい程マクロ組織の微細化をもたらす。ストリップキャスティング法はこの効果により、α−Ｆｅ生成を抑制し、良好なＲの含有量が多い相３の分散に成功している。しかし、Ｒ₂ Ｆｅ₁₄Ｂ相結晶粒１の微細化によって、微粉砕後の粉砕粉中に方位の異なる２つ以上のＲ₂ Ｆｅ₁₄Ｂ相１が存在する確率が高まり、配向性の低下をもたらすため、冷却速度を適度に制御することが必要となる。本発明の合金に於ても８００℃以上、特に包晶反応終了までの冷却速度を主相の短軸方向の大きさ（Ｗ）が、１０μｍ未満にならない範囲内で適度に増加させることは、共晶領域２の分散性向上によるＲの含有量が多い相３の分散性の向上をもたらし有効である。
【００２２】
【作用】
本発明はＲ（希土類元素）を含む永久磁石の原料合金に於て、従来単に粒界相又はＲリッチ相と呼ばれていた、鋳造時に最後まで液相として存在するＲに富んだ部分が凝固した領域（共晶領域）内の組織を制御することにより、実質的に利用できるＲに富んだ相の体積率を増加すること、さらに主相の結晶粒径、共晶領域の分散状態の制御により、高特性Ｒ−Ｆｅ−Ｂ系焼結磁石用の原料として適した合金を提供するものである。特に従来注目されていなかった包晶反応によって主相Ｒ₂ Ｆｅ₁₄Ｂ相が生成した後の、８００℃から６００℃での凝固冷却速度に注目し、複雑な工程、装置を用いることなく優れた焼結磁石用合金を提供するものである。
【００２３】
【実施例】
以下、実施例により本発明を更に詳細に説明する。
（実施例１）
表１に示すように、合金インゴットの組成が、Ｎｄ：２９．０重量％、Ｄｙ：３．４重量％、Ｂ：１. ０重量％、Ａｌ：０. ３５重量％、残部鉄になるように、鉄ネオジム合金、金属ディスプロシウム、フェロボロン、アルミニウム、鉄を配合し、アルゴンガス雰囲気中で、アルミナるつぼを使用して高周波溶解炉で溶解し、銅製箱型鋳型に鋳造した。この際、鋳型に設置した熱電対で合金の凝固時の温度変化を測定し、８００℃〜６００℃での冷却速度は表１に示すように平均１７℃／秒であった。なお、モールド比は２０であり、得られたインゴットの厚さは５ｍｍであった。その断面のマクロ組織は鋳型近傍のチル晶部を除いて柱状晶であり、さらに研磨した後、その断面の組織を偏光顕微鏡で観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の粒径は長軸方向２００〜１０００μｍ程度、短軸方向２０〜１００μｍ程度であり、柱状晶部分のインゴット全体に対する体積率は８５％であった。また、反射電子顕微鏡で観察した結果、共晶領域は主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の長軸方向に細長く伸長し、その間隔は１０〜３０μｍ程度であり、共晶領域内に析出するＲの含有量がより少ない相の短径は１μｍ程度であった。
【００２４】
次に得られた合金インゴットを、窒素ガス中においてブラウンミルで３５メッシュ以下まで粉砕した後、さらに窒素ガス中においてジェットミルで４μｍまで微粉砕した。次いで得られた微粉末を１０ＫＯｅ、１ｔｏｎｆ／ｃｍ² の条件で磁場成形し、１０×１０×１０ｍｍの成形体を得た後、真空中１０６０℃にて２時間焼結し、さらに真空中６２０℃にて１時間の時効処理した。得られた焼結磁石の磁気特性を表１に合せて示す。最大磁気エネルギー積は３９．８ＭＧＯｅ、保磁力は２０．０ｋＯｅであった。
【００２５】
（実施例２）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、銅製水冷箱型鋳型に鋳造した。この際、実施例１と同様に凝固時の温度変化を測定し、８００℃〜６００℃での冷却速度は平均１２℃／秒であった。鋳型の水冷は鋳型外周部にろう付けした銅管内に水を流した。なお、モールド比は１０であり、得られたインゴットの厚さは５ｍｍであった。その断面のマクロ組織は鋳型近傍のチル晶部を除いて柱状晶であり、さらに研磨した後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向２００〜１０００μｍ程度、短軸方向２０〜１００μｍ程度であった。柱状晶部分のインゴット全体に対する体積率は８３％であった。また、共晶領域は実施例１で作製したインゴットと同様に主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の長軸方向に細長く伸長し、その間隔は１０〜４０μｍ程度であり、共晶領域内に析出するＲの含有量がより少ない相の短径は２μｍ程度であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３８．８ＭＧＯｅ、保磁力は１９．３ｋＯｅであった。
【００２６】
（実施例３）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、鉄製水冷箱型鋳型に鋳造した。この際、実施例１と同様に凝固時の温度変化を測定し、８００℃〜６００℃での冷却速度は平均８℃／秒であった。鋳型の水冷は鋳型側板の厚み方向中央部に設けた直径２０ｍｍの穴に水を流した。なお、モールド比は２０であり、得られたインゴットの厚さは１０ｍｍでった。その断面のマクロ組織は鋳型近傍のチル晶部を除いて柱状晶であり、さらに研磨した後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向３００〜１５００μｍ、短軸方向３０〜１５０μｍであった。柱状晶部分のインゴット全体に対する体積率は９２％であった。また、共晶領域は実施例１で作製したインゴットと同様に主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の長軸方向に細長く伸長し、その間隔は１０〜５０μｍ程度であり、共晶領域内に析出するＲの含有量がより少ない相の短径は２μｍ程度であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３８．５ＭＧＯｅ、保磁力は１９．１ｋＯｅであった。
【００２７】
（実施例４）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、鋳型内径５００ｍｍ長さ１０００ｍｍの遠心鋳造装置にて鋳造した。この時の鋳型の回転数は、遠心力が１０Ｇとなるように、１８９ｒｐｍに設定し、溶湯供給終了後に鋳型内部にアルゴンガスを供給してインゴットを冷却した。なお、ここでは凝固時の温度測定は実施していない。得られた合金インゴットの厚さは２〜３ｍｍであり、その断面のマクロ組織は鋳型近傍のチル晶部を除いて柱状晶であった。その後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向３００〜１０００μｍ、短軸方向３０〜１００μｍであった。柱状晶部分のインゴット全体に対する体積率は９８％であった。また、共晶領域は実施例１で作製したインゴットと同様に主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の長軸方向に細長く伸長し、その間隔は１０〜３０μｍ程度であり、共晶領域内に析出するＲの含有量がより少ない相の短径は１μｍ程度であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３９．５ＭＧＯｅ、保磁力は１９．５ｋＯｅであった。
【００２８】
（比較例１）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、銅製箱型鋳型に鋳造した。この際、実施例１と同様に凝固時の温度変化を測定し、８００℃〜６００℃での冷却速度は平均３℃／秒であった。なおモールド比は５であり、得られたインゴットの厚さは５ｍｍであった。その断面のマクロ組織は鋳型近傍のチル晶部を除いて柱状晶であり、さらに研磨した後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向２００〜１０００μｍ程度、短軸方向２０〜１００μｍ程度であった。柱状晶部分のインゴット全体に対する体積率は８１％であった。また、共晶領域は実施例１で作製したインゴットと同様に主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の長軸方向に細長く伸長し、その間隔は１０〜４０μｍ程度であり、共晶領域内に析出するＲの含有量がより少ない相の短径は５μｍ程度であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３６．５ＭＧＯｅ、保磁力は１７．０ｋＯｅであった。
【００２９】
（比較例２）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、鉄製水冷箱型鋳型に鋳造した。この際、実施例１と同様に凝固時の温度変化を測定し、８００℃〜６００℃での冷却速度は平均１℃／秒であった。なお、鋳型の水冷は実施例２と同様の機構とし、モールド比は５であり、得られたインゴットの厚さは２０ｍｍであった。その断面のマクロ組織は殆ど柱状晶であるが、鋳型近傍にチル晶組織、インゴット中央部には等軸晶組織が存在していた。さらに研磨した後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向３００〜１５００μｍ程度、短軸方向３０〜２００μｍ程度であった。柱状晶部分のインゴット全体に対する体積率は８０％であった。また、共晶領域は主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の主に粒界部に存在し、特に等軸晶領域では粒界三重点で５０μｍ程度まで粗大化していた。また、共晶領域内に析出するＲの含有量がより少ない相の短径は５μｍ以上であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３５．２ＭＧＯｅ、保磁力は１５．５ｋＯｅであった。
【００３０】
（実施例５）
比較例２で作製したインゴットを直径５ｍｍ以下に割り、アルゴン雰囲気中で９００℃、１時間の熱処理を実施した後、アルゴンガスを直接吹き付け冷却した。その際の試料表面の温度変化を測定したところ、８００℃〜６００℃での冷却速度は平均１０℃／秒であった。該合金のマクロ組織、主相の粒径、共晶領域の分散状態に変化はなかったが、共晶領域内に析出するＲの含有量がより少ない相の短径は３μｍ程度であった。次に得られた試料より、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３７．５ＭＧＯｅ、保磁力は１８．０ｋＯｅであった。
【００３１】
（実施例６）
実施例１と同じ組成となるように、実施例１と同じ方法で溶解し、直径５０ｃｍの銅製ロールを用いて単ロール型ストリップキャスティング法により合金箔片を作製した。なお、ここでは凝固時の温度測定は実施していない。この時のロール周速度は３ｍ／ｓであった。得られる合金箔片の厚さは０．１〜０．２ｍｍと極く薄くしてある。その断面のマクロ組織はロール接触面側がチル晶であり、他は柱状晶であった。さらに研磨した後、実施例１と同様の方法で組織観察した結果、柱状晶部の主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の平均粒径は長軸方向３０〜５０μｍ程度、短軸方向２〜８μｍ程度であった。柱状晶部分のインゴット全体に対する体積率は７０％であった。また、共晶領域は主相Ｎｄ₂ Ｆｅ₁₄Ｂ相の主に粒界部に存在しており、共晶領域内に析出するＲの含有量がより少ない相の短径は１μｍ以下であった。次に得られたインゴットより、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３７．１ＭＧＯｅ、保磁力は２０．１ｋＯｅであった。
【００３２】
（比較例３）
比較例２で作製したインゴットを厚さ２０ｍｍのままアルゴン雰囲気中で９００℃、１時間の熱処理を実施した後、アルゴンガスを直接吹き付け冷却した。その際の試料の温度変化を測定したところ、８００℃〜６００℃での冷却速度は平均３℃／秒であった。該インゴットのマクロ組織、主相の粒径、共晶領域の分散状態に変化はなく、共晶領域内に析出するＲの含有量がより少ない相の粒径も熱処理前同様に５μｍ以上であった。次に得られた試料より、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３４．９ＭＧＯｅ、保磁力は１５．１ｋＯｅであった。
【００３３】
（比較例４）
比較例２で作製したインゴットを厚さ２０ｍｍのままアルゴン雰囲気中で６００℃、１時間の熱処理を実施した後、アルゴンガスを直接吹き付け冷却した。該インゴットのマクロ組織、主相の粒径、共晶領域の分散状態に変化はなく、共晶領域内に析出するＲの含有量がより少ない相の粒径も熱処理前同様に５μｍ以上であった。次に得られた試料より、実施例１と同様の方法で焼結磁石を作製し、その最大エネルギー積は３５．１ＭＧＯｅ、保磁力は１５．３ｋＯｅであった。
【００３４】
【表１】

【００３５】
【発明の効果】
本発明によれば、高性能希土類磁石用原料として最適な原料合金を複雑な工程、装置を用いることなく製造することが可能となり、極めて有用である。
【図面の簡単な説明】
【図１】本発明による合金の顕微鏡組織を模式的に示す図である。
【図２】従来の合金の顕微鏡組織を模式的に示す図である。
【図３】本発明による合金の共晶領域の組織を拡大して示した模式図である。
【図４】従来の合金の共晶領域の組織を拡大して示した模式図である。
【符号の説明】
１Ｒ₂ Ｆｅ₁₄Ｂ柱状晶
２共晶領域
３Ｒの含有量がより多い相
４Ｒの含有量がより少ない相[0001]
[Industrial application fields]
The present invention relates to a raw material alloy used as a raw material for a magnet containing a rare earth element and a method for producing the raw material alloy.
[0002]
[Prior art]
Recently, rare earth sintered magnets or rare earth bonded magnets that make use of the excellent magnetic properties of rare earth alloys have attracted attention. In particular, R-Fe-B magnets have been developed with further improved magnetic properties. It has been broken. In R-Fe-B magnets, the ferromagnetic phase R responsible for magnetism ₂ Fe ₁₄ In addition to the B phase, there is a non-magnetic phase with a high concentration of rare earth elements such as Nd (referred to as an R-rich phase), which plays an important role as follows.
{Circle around (1)} The melting point is low and it becomes a liquid phase at the time of sintering in the magnetizing process, which contributes to increasing the density of the magnet and thus improving the magnetization.
(2) Eliminate grain boundary irregularities, reduce the reverse creation of new domains and increase coercivity.
(3) The R-rich phase is non-magnetic and magnetically insulates the main phase, thus increasing the coercive force.
Therefore, if there is an interface that is not covered by the R-rich phase because the volume ratio of the R-rich phase is low or the dispersion state is bad, the squareness deteriorates due to the local coercive force reduction in that portion, It is known that the magnetization is also reduced due to poor sintering, so that the maximum magnetic energy product is reduced.
[0003]
However, the higher the characteristic magnet, the more the R phase is ferromagnetic. ₂ Fe ₁₄ Since it is necessary to increase the volume ratio of the B phase, the volume ratio of the R-rich phase inevitably decreases, resulting in a partial shortage of the R-rich phase, and sufficient characteristics are often not obtained. Therefore, many studies have been reported regarding prevention of deterioration of properties due to lack of R-rich phase of high-quality materials, and they are roughly divided into two groups.
[0004]
One is the main phase R ₂ Fe ₁₄ The B phase and the R rich phase are supplied from different alloys, and is generally called a two alloy method. Almost R ₂ Fe ₁₄ In order to improve the dispersibility of the R-rich phase, two types of alloys, a B-phase single phase and an alloy that generates an R-rich phase, are each finely pulverized and then mixed and molded and sintered at an appropriate ratio. Many innovations are found in alloys that produce rich phases. For example, if an amorphous alloy having a liquid phase composition at the sintering temperature is used, the content of Fe is higher than that of the R-rich phase. The mixing ratio with the alloy that generates the main phase can be increased. As a result, the dispersibility of the R-rich phase generated during sintering is improved, and the magnetic properties have been improved (E. Otsuki, T. Otsuka and T.). Imai, 11th Internatinal Workshop on Rare Earth magnets and their Applications, vol.1, p328 (1990)).
[0005]
The other is to produce an alloy having a structure in which the R-rich phase is finely dispersed by solidification at a higher cooling rate than in the conventional mold casting method by the strip casting method. Since the R-rich phase in the alloy is finely dispersed, the dispersibility of the R-rich phase after pulverization and sintering is improved, and the magnetic properties have been successfully improved (Japanese Patent Laid-Open No. 5-222488, Japanese Patent Laid-Open No. 295490).
[0006]
[Problems to be solved by the invention]
As described above, the two-alloy method and the strip casting method have resulted in good dispersion of the R-rich phase and improved magnetic properties. However, these methods have the following problems.
[0007]
First, in the former two-alloy method, extremely high properties have been reported. However, liquid ultra-quenching must be used to produce an amorphous alloy that generates an R-rich phase. It will cost worse. There is also a report that uses Co-based intermetallic compounds in R-rich phase-forming alloys without using liquid superquenching, and has developed characteristics of 45 MGOe or more (Tatsumi, Minowa, Motoshima, IEEJ Transaction A, Vol. 113, Volume 12) No., P849-853, 1993). But R ₂ Fe ₁₄ Since the B phase is generated from α-Fe and the liquid phase by the peritectic reaction, it is almost R ₂ Fe ₁₄ An alloy that generates a main phase that requires casting with a B stoichiometric composition has a very large residual amount of α-Fe. Α-Fe remarkably impairs pulverizability, causes composition fluctuations during pulverization, and causes a decrease in magnetic properties and an increase in variation. Therefore, the alloy that generates the main phase eliminates α-Fe that is the primary crystal, and almost R ₂ Fe ₁₄ A homogenization process for a long time is required from the necessity of using the B phase as a single phase. Furthermore, in the two alloy method, two types of alloy powders must be pulverized separately and mixed uniformly, so the productivity is much worse and the cost increases compared to the conventional method of pulverizing one alloy of as cast. Since it cannot be avoided, it can be used only for extremely high-quality materials with high added value of about 45MGOe or more.
[0008]
On the other hand, the strip casting method also takes a long time to cast, and since the molten metal containing a highly active rare earth element is held for a long time and supplied little by little, the components are likely to fluctuate due to the reaction between the crucible, holding furnace or tundish and molten metal. . In addition, it is extremely difficult to maintain a constant temperature and maintain stable casting in a steady state, and there is a problem that the yield is low. In addition, there is a problem that a special and expensive casting facility is required, and an increase in cost is inevitable.
[0009]
[Means for Solving the Problems]
As a result of studying a method for producing an alloy that is unlikely to cause characteristic deterioration due to lack of R-rich phase and uneven distribution by a conventional mold casting method without using a special and expensive casting equipment, the present inventor has conventionally simply used a grain boundary. Attention was paid to a region (referred to as eutectic region), which was called a phase or R-rich phase, which solidified an R-rich portion that existed as a liquid phase until the end of casting. As a result, we found the fact that the morphology of the eutectic part changes depending on the cooling rate around 600-800 ° C, and that it is related to the coercivity and squareness of the magnet. Furthermore, it was found that the effect is further enhanced by controlling the crystal grain size of the main phase and the dispersion state of the eutectic region. The present invention has been made based on these findings.
[0010]
That is, the present invention is a conventional raw material alloy for permanent magnets based on R (at least one of rare earth elements including Y), T (transition metal in which Fe is essential) and B (boron). By controlling the structure of the eutectic region, where the R-rich portion, which was called the boundary phase or R-rich phase, solidified as a liquid phase until the end of casting, is substantially usable, This problem is solved by increasing the volume fraction of the phase and controlling the crystal grain size of the main phase and the dispersion state of the eutectic region.
[0011]
Next, the configuration of the present invention will be described in detail below.
1 and 2 are schematic views of typical microstructures of the present invention and conventional alloys. In the conventional alloy shown in FIG. 2, the main phase 1 crystallizes in a columnar shape and exhibits a structure surrounded by a eutectic region 2 that solidifies last. Further, in the alloy of the present invention shown in FIG. 1, the eutectic region 2 is also present inside the crystal grains of the main phase 1. 2 is somewhat enlarged, and the size of the main phase 1 is larger in the alloy of the present invention in FIG.
(1) Eutectic region structure
The minor axis direction of the phase having a smaller R content in the eutectic region is 3 μm or less.
The portion rich in R (rare earth) that exists as a liquid phase until the end is finally transformed and solidified into several solid phases between 600 and 800 ° C. by eutectic reaction. For example, in the Nd—Fe—B ternary system, the liquid phase is an Nd metal phase, Nd by an eutectic reaction at 665 ° C. ₂ Fe ₁₄ B phase, NdFe _Four B _Four It is known to transform and solidify into three phases. When other component elements are included, the reaction temperature and product phase are considered to change. Originally, the Nd-rich phase refers to the Nd metal phase in the Nd—Fe—B ternary system, but Nd produced by the eutectic reaction. ₂ Fe ₁₄ B phase, NdFe _Four B _Four There are cases in which the phases are treated as a collective term, especially when discussing the structure of the raw material alloy. The same applies when other component elements are included. Therefore, in the present invention, the region where the R-rich portion, which has been widely known as the R-phase or grain boundary phase in the past, is solidified as the liquid phase during casting is called the eutectic region 2 and is the primary solid solution of R. This is expressed separately from the R-rich phase.
[0012]
FIG. 3 shows an enlarged schematic view of the structure in the eutectic region of the alloy according to the present invention. The eutectic region 2 has a structure in which the phase 4 having a smaller R content than the surroundings is precipitated in a thin rod shape. The phase 4 having a lower R content in the eutectic region 2 is the R shown above. ₂ Fe ₁₄ Phase B, RFe _Four B _Four Although it is a phase and a different phase may exist when it contains other component elements, it can be said that it is a phase other than the surrounding phase 3 having a large R content. If the size in the minor axis direction of the phase 4 having a smaller R content is about the finely pulverized particle size for obtaining a magnetic field forming powder in the magnetizing step, that is, 3 to 5 μm or less, these phases 4 are The pulverized powder contained always coexists with the phase 3 having a high R content. Therefore, the ratio of the powder containing the phase 3 having a large R content increases, the dispersibility of the phase 3 having a large R content is improved, the densification during the sintering is promoted, and the heat treatment after the sintering is further promoted. The temperature dependency of the coercive force is improved, and a magnet having excellent magnetic properties can be manufactured.
[0013]
On the other hand, FIG. 4 shows an enlarged schematic view of the structure in the eutectic region of the alloy according to the conventional method. In the conventional alloys, the size of the phase 4 having a smaller R content is thicker than in the case of the present invention, and some of them have a flake shape. Compared to the alloy according to the present invention, the phase 4 having a smaller R content inside the eutectic region 2 is larger in the minor axis direction than the finely pulverized particle size and includes the phase 3 having a larger R content. Since the proportion of the powder that does not increase increases, the dispersibility of phase 3 having a high R content decreases, and the magnetic properties decrease. In order to refine the phase 4 having a smaller R content, it is effective to increase the cooling rate during the eutectic reaction for producing these phases, and to prevent the coarsening of each produced phase while increasing the nucleation frequency. is there. There is also a possibility that miniaturization is promoted by a certain kind of additive element.
[0014]
(2) Main phase crystal grain size
R is the main phase ₂ Fe ₁₄ The size of the B crystal grains in the major axis direction is 50 μm or more, the size in the minor axis direction is 10 μm or more, and the volume ratio of this main phase region is 70% or more.
In the microscopic structure shown in FIG. 1 cut along the solidification direction, in one pulverized powder, the crystal grain size of the main phase 1 is about the finely pulverized particle size for obtaining a magnetic field forming powder, that is, about 3 to 5 μm or less. In this case, two or more main phases having different orientations exist, and the orientation deteriorates. Therefore, it is convenient that the crystal grain size of the main phase 1 is large, and the volume ratio of crystal grains having a major axis size (L) of 50 μm or more and a minor axis size (W) of 10 μm or more is 70. % Or more is preferable. More preferably, the size (L) in the major axis direction is 100 μm or more, and the size (W) in the minor axis direction is 20 μm or more. Each crystal grain of the main phase can be easily identified by observing the surface buffed using alumina, diamond or the like after polishing the alloy with emery paper. In the polarizing microscope, incident polarized light is reflected by rotation of the polarization plane corresponding to the magnetization direction of the ferromagnetic surface due to the magnetic Kerr effect, so that the difference in the polarization plane reflected from each crystal grain is observed as light and dark.
Looking at the polarization micrograph of the alloy of the present invention, the main phase R ₂ Fe ₁₄ B grain boundaries are observed. The difference in crystal direction of each crystal grain becomes bright and dark and can be clearly identified. Further, the eutectic region observed as black streaks is not only the main phase grain boundary but also the main phase R ₂ Fe ₁₄ The B crystal grains also exist almost in parallel with the major axis direction. The fine stripe pattern in the main phase crystal grain corresponds to the magnetic domain.
[0015]
(3) Distribution of eutectic region
Eutectic region 2 is the main phase R ₂ Fe ₁₄ In addition to the grain boundary of B1, it also exists in the crystal grains of the main phase 1. In particular, the interval (D) between the eutectic regions 2 is the main phase R. ₂ Fe ₁₄ The volume ratio of the region which is less than or equal to one half of the size (W) of the B crystal grain 1 in the minor axis direction is 70% or more.
Here, the interval (D) between the eutectic regions 2 may be the eutectic regions in the crystal grain boundaries or the eutectic regions in the crystal grains. Alternatively, it may be the interval between the crystal grain boundary and the eutectic region in the crystal grain. That is, the interval between adjacent eutectic regions may be less than or equal to one half of the size of the main phase in the minor axis direction.
When the alloy is finely pulverized in the item (1), the ratio of the powder containing the phase 3 having a high R content increases due to the refinement of the structure inside the eutectic region 2, and the phase 3 having a high R content. It was noted that the improvement of dispersibility makes it possible to improve the sinterability and magnetic properties. The dispersibility of the phase 3 having a high R content is further improved by finely dispersing the eutectic region 2 including the phase 3 having a high R content in the alloy. That is, according to the expansion of the interface between the eutectic region 2 and the main phase 1, the amount of pulverized powder containing the phase 3 having a large R content during pulverization increases, and therefore the phase 3 having a high R content. Improves dispersibility. Therefore, it is convenient that the eutectic region 2 is finely dispersed. Specifically, the interval (D) between the eutectic regions 2 is R of the main phase. ₂ Fe ₁₄ It is preferable that the volume ratio of the region which is less than or equal to one half of the size (W) in the minor axis direction of the B crystal grains 1 is 70% or more.
[0016]
In the alloy of the present invention, the eutectic region 2 is a main phase R in a columnar crystal structure. ₂ Fe ₁₄ Some of the B1 crystal grain boundaries and crystal grains extend parallel to the major axis direction, and some equiaxed crystal structures exist in a spherical shape. Therefore, in order to expand the interface between the eutectic region 2 and the main phase 1 and increase the dispersibility of the phase 3 having a large R content, a columnar crystal structure in which the eutectic region 2 is thin and elongated is preferable. On the other hand, the alloy obtained by the conventional casting method includes the main component R constituting the crystal grains including the strip casting method. ₂ Fe ₁₄ B1 is a structural structure (Japanese Patent Laid-Open No. 5-295490) covered with phase 3 having a large content of R constituting the grain boundary, and main phase R according to the present invention ₂ Fe ₁₄ There is no description about the existence of the eutectic region 2 including the phase 3 having a large content of a plurality of R in the crystal grains of B1, and it can be said that it is a completely different structure.
[0017]
(4) Manufacturing method
The method for producing the alloy of the present invention will be described. One is to heat the alloy ingot to 800 ° C. to 1150 ° C. after casting and cool the temperature range from 800 ° C. to 600 ° C. at a cooling rate of 5 ° C./second or more. Characterized by
The other is to cool a temperature range from 800 ° C. to 600 ° C. at a cooling rate of 5 ° C./second or more during casting.
Hereinafter, each process will be described.
[0018]
The structure in which the size (W) in the minor axis direction of the phase 4 having a smaller R content in the eutectic region 2 according to the present invention is 3 μm or less is generated by heat treatment after casting. That is, after heating the solidified alloy once and remelting the eutectic region 2, the cooling rate in the temperature range where the eutectic region 2 solidifies is increased, and the nucleation frequency is increased and each generated phase is coarsened. By preventing, the phase 4 with less R content in the eutectic region 2 is refined. Although the solidification temperature varies somewhat depending on the composition, the specific heating temperature is 800 ° C. or higher, which is higher than the temperature at which the eutectic region 2 melts, and the main phase R ₂ Fe ₁₄ It is necessary to set the temperature to 1150 ° C. or lower at which B phase 1 melts. The cooling is preferably performed in a temperature range of 800 ° C. to 600 ° C. at a cooling rate of 5 ° C./second or more, and a more preferable cooling rate is 10 ° C./second or more. In order to increase the cooling rate, gas quenching that does not require a special device is possible, but in order to increase the cooling efficiency, it is preferable to finely pulverize the alloy ingot to about 5 mm.
[0019]
Next, a method for controlling the cooling rate during casting will be described.
Increasing the cooling rate in the temperature range where the eutectic region 2 is formed, that is, the temperature range where the R (rare earth) -rich portion that exists as a liquid phase until the end solidifies, increases the nucleation frequency and increases the coarseness of each generated phase By preventing the formation, the phase 4 having a smaller R content in the eutectic region 2 is refined. Depending on the composition, the solidification temperature slightly rises and falls. Specifically, it is preferable to cool a temperature range of 800 ° C. to 600 ° C. at a cooling rate of 5 ° C./second or more, and a more preferable cooling rate is 10 ° C./second or more. . Such a condition can be achieved without using a special casting method such as a strip casting method, for example, even by a conventional casting method by sufficiently increasing the mold ratio (mold ratio).
[0020]
The mold ratio is a value obtained by dividing the heat capacity of the mold by the heat capacity of the alloy to be cast, and can be easily compared by simply comparing the respective weight ratios and, more simply, the ratio of the thicknesses of the respective side plates. For example, since iron and copper, which are representative mold materials, have substantially the same heat capacity per unit volume, it is possible to compare them including differences in materials when discussed in terms of the side plate thickness ratio. Since the mold temperature increase is suppressed by increasing the mold ratio, it is possible to increase the cooling rate at 800 ° C. or lower where the cooling efficiency of the ingot has been particularly lowered due to the temperature increase of the mold. The mold ratio can be increased by increasing the thickness of the side plate or increasing the thickness of the ingot. Furthermore, a method of quickly removing the heat transferred from the molten metal by having an efficient water cooling mechanism and preventing the temperature of the mold from rising is also effective, and a combination of both enables more efficient cooling. In addition, it is also effective to make ingots with a centrifugal casting method.
[0021]
In the present invention, a cooling rate of 800 ° C. or higher and 600 ° C. or lower is not particularly specified. ₂ Fe ₁₄ It is well known that the crystal grain size of the B phase 1 and the dispersibility of the eutectic region 2 are affected. In a general R—Fe—B magnet composition, primary α-Fe is formed from the liquid phase, and R and R are formed by peritectic reaction from the liquid phase. ₂ Fe ₁₄ When B phase 1 is formed or when the solidification rate is fast, it is subcooled to a peritectic reaction temperature or lower and R is removed from the liquid phase. ₂ Fe ₁₄ B phase 1 is produced directly. Therefore, the influence of the cooling rate until the peritectic reaction on the macroscopic alloy structure is particularly great, and the macrostructure becomes finer as the cooling rate increases. Due to this effect, the strip casting method suppresses the formation of α-Fe and succeeds in dispersing the phase 3 having a good R content. But R ₂ Fe ₁₄ Two or more Rs having different orientations in the pulverized powder after pulverization due to the refinement of the B phase crystal grains 1 ₂ Fe ₁₄ Since the probability that the B phase 1 exists increases and the orientation is lowered, it is necessary to appropriately control the cooling rate. Even in the alloy of the present invention, the cooling rate until 800 ° C. or more, particularly the peritectic reaction, is increased appropriately within the range where the minor axis size (W) of the main phase is not less than 10 μm. This is effective because the dispersibility of the phase 3 having a large R content is improved by improving the dispersibility of the eutectic region 2.
[0022]
[Action]
In the present invention, a permanent magnet raw material alloy containing R (rare earth element), which has been conventionally called simply a grain boundary phase or an R-rich phase, the R-rich portion existing as a liquid phase at the end of casting is solidified. By controlling the structure in the region (eutectic region), the volume fraction of the R-rich phase that can be substantially used is increased, and the crystal grain size of the main phase and the dispersion state of the eutectic region are controlled. Thus, an alloy suitable as a raw material for a high-characteristic R—Fe—B based sintered magnet is provided. In particular, the main phase R is due to the peritectic reaction, which has not attracted much attention in the past. ₂ Fe ₁₄ Focusing on the solidification cooling rate at 800 ° C. to 600 ° C. after the generation of the B phase, the present invention provides an excellent alloy for sintered magnets without using complicated processes and apparatuses.
[0023]
【Example】
Hereinafter, the present invention will be described in more detail with reference to examples.
(Example 1)
As shown in Table 1, the composition of the alloy ingot is Nd: 29.0 wt%, Dy: 3.4 wt%, B: 1.0 wt%, Al: 0.35 wt%, and the balance iron. In addition, an iron neodymium alloy, metal dysprosium, ferroboron, aluminum, and iron were blended, melted in a high-frequency melting furnace using an alumina crucible in an argon gas atmosphere, and cast into a copper box mold. At this time, the temperature change at the time of solidification of the alloy was measured with a thermocouple installed in the mold, and the cooling rate at 800 ° C. to 600 ° C. was an average of 17 ° C./second as shown in Table 1. The mold ratio was 20, and the thickness of the obtained ingot was 5 mm. The macro structure of the cross section is a columnar crystal excluding the chill crystal part in the vicinity of the mold, and after further polishing, the structure of the cross section is observed with a polarizing microscope. As a result, the main phase Nd of the columnar crystal part is observed. ₂ Fe ₁₄ The particle size of the B phase was about 200 to 1000 μm in the major axis direction and about 20 to 100 μm in the minor axis direction, and the volume ratio of the columnar crystal portion to the entire ingot was 85%. Further, as a result of observation with a reflection electron microscope, the eutectic region was found to have a main phase Nd. ₂ Fe ₁₄ The phase of the phase B was elongated in the major axis direction, the interval was about 10 to 30 μm, and the minor axis of the phase with less R content precipitated in the eutectic region was about 1 μm.
[0024]
Next, the obtained alloy ingot was pulverized in a nitrogen gas to 35 mesh or less in a nitrogen gas and further pulverized in a nitrogen gas to 4 μm by a jet mill. Next, the fine powder obtained was 10 KOe, 1 tonf / cm. ² After forming a 10 × 10 × 10 mm molded body under the above conditions, it was sintered in vacuum at 1060 ° C. for 2 hours and further subjected to aging treatment in vacuum at 620 ° C. for 1 hour. Table 1 shows the magnetic properties of the obtained sintered magnet. The maximum magnetic energy product was 39.8 MGOe, and the coercive force was 20.0 kOe.
[0025]
(Example 2)
It melt | dissolved by the same method as Example 1 so that it might become the same composition as Example 1, and it casted to the copper water-cooled box type | mold mold. At this time, the temperature change during solidification was measured in the same manner as in Example 1, and the cooling rate at 800 ° C. to 600 ° C. was an average of 12 ° C./second. In the water cooling of the mold, water was poured into a copper tube brazed to the outer periphery of the mold. The mold ratio was 10, and the thickness of the obtained ingot was 5 mm. The macrostructure of the cross section is a columnar crystal except for the chill crystal part in the vicinity of the mold, and after further polishing, the structure was observed by the same method as in Example 1. As a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was about 200 to 1000 μm in the major axis direction and about 20 to 100 μm in the minor axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 83%. Further, the eutectic region is the main phase Nd as in the ingot produced in Example 1. ₂ Fe ₁₄ The phase of the phase B was elongated in the major axis direction, the interval was about 10 to 40 μm, and the minor axis of the phase with less R content precipitated in the eutectic region was about 2 μm. Next, a sintered magnet was produced from the obtained ingot in the same manner as in Example 1, and the maximum energy product was 38.8 MGOe and the coercive force was 19.3 kOe.
[0026]
(Example 3)
It melt | dissolved by the same method as Example 1 so that it might become the same composition as Example 1, and it casted to the iron water-cooled box type | mold mold. Under the present circumstances, the temperature change at the time of solidification was measured like Example 1, and the cooling rate in 800 to 600 degreeC was an average of 8 degree-C / sec. In the water cooling of the mold, water was poured into a hole having a diameter of 20 mm provided in the central portion in the thickness direction of the mold side plate. The mold ratio was 20, and the thickness of the obtained ingot was 10 mm. The macrostructure of the cross section is a columnar crystal except for the chill crystal part in the vicinity of the mold, and after further polishing, the structure was observed by the same method as in Example 1. As a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was 300 to 1500 μm in the long axis direction and 30 to 150 μm in the short axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 92%. Further, the eutectic region is the main phase Nd as in the ingot produced in Example 1. ₂ Fe ₁₄ The phase of the phase B was elongated in the major axis direction, the interval was about 10 to 50 μm, and the minor axis of the phase with less R content precipitated in the eutectic region was about 2 μm. Next, from the obtained ingot, a sintered magnet was produced in the same manner as in Example 1, and the maximum energy product was 38.5 MGOe and the coercive force was 19.1 kOe.
[0027]
(Example 4)
It melt | dissolved by the same method as Example 1 so that it might become the same composition as Example 1, and it casted with the centrifugal casting apparatus of mold internal diameter 500mm length 1000mm. The rotational speed of the mold at this time was set to 189 rpm so that the centrifugal force was 10 G, and after the molten metal supply was completed, argon gas was supplied into the mold to cool the ingot. In addition, the temperature measurement at the time of solidification is not implemented here. The thickness of the obtained alloy ingot was 2 to 3 mm, and the macrostructure of the cross section was a columnar crystal except for the chill crystal part in the vicinity of the mold. Thereafter, the structure was observed in the same manner as in Example 1, and as a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was 300 to 1000 μm in the major axis direction and 30 to 100 μm in the minor axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 98%. Further, the eutectic region is the main phase Nd as in the ingot produced in Example 1. ₂ Fe ₁₄ The phase of the phase B was elongated in the major axis direction, the interval was about 10 to 30 μm, and the minor axis of the phase with less R content precipitated in the eutectic region was about 1 μm. Next, from the obtained ingot, a sintered magnet was produced by the same method as in Example 1, and the maximum energy product was 39.5 MGOe and the coercive force was 19.5 kOe.
[0028]
(Comparative Example 1)
It melt | dissolved by the same method as Example 1 so that it might become the same composition as Example 1, and it casted to the copper box type | mold mold. At this time, the temperature change during solidification was measured in the same manner as in Example 1, and the cooling rate at 800 ° C. to 600 ° C. was an average of 3 ° C./second. The mold ratio was 5, and the thickness of the obtained ingot was 5 mm. The macrostructure of the cross section is a columnar crystal except for the chill crystal part in the vicinity of the mold, and after further polishing, the structure was observed by the same method as in Example 1. As a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was about 200 to 1000 μm in the major axis direction and about 20 to 100 μm in the minor axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 81%. Further, the eutectic region is the main phase Nd as in the ingot produced in Example 1. ₂ Fe ₁₄ The phase of the phase B was elongated in the major axis direction, the interval was about 10 to 40 μm, and the minor axis of the phase with less R content precipitated in the eutectic region was about 5 μm. Next, from the obtained ingot, a sintered magnet was produced in the same manner as in Example 1. The maximum energy product was 36.5 MGOe and the coercive force was 17.0 kOe.
[0029]
(Comparative Example 2)
It melt | dissolved by the same method as Example 1 so that it might become the same composition as Example 1, and it casted to the iron water-cooled box type | mold mold. Under the present circumstances, the temperature change at the time of solidification was measured like Example 1, and the cooling rate in 800 to 600 degreeC was an average of 1 degree-C / sec. The mold was cooled by water in the same manner as in Example 2, the mold ratio was 5, and the thickness of the obtained ingot was 20 mm. The macro structure of the cross section was almost columnar, but a chill crystal structure was present near the mold and an equiaxed crystal structure was present in the center of the ingot. After further polishing, the structure was observed by the same method as in Example 1. As a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was about 300 to 1500 μm in the major axis direction and about 30 to 200 μm in the minor axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 80%. The eutectic region is the main phase Nd. ₂ Fe ₁₄ The B phase mainly exists in the grain boundary part, and in the equiaxed crystal region, the grain boundary triple point is coarsened to about 50 μm. In addition, the minor axis of the phase with less R content precipitated in the eutectic region was 5 μm or more. Next, from the obtained ingot, a sintered magnet was produced in the same manner as in Example 1. The maximum energy product was 35.2 MGOe and the coercive force was 15.5 kOe.
[0030]
(Example 5)
The ingot produced in Comparative Example 2 was divided into a diameter of 5 mm or less, and after heat treatment at 900 ° C. for 1 hour in an argon atmosphere, argon gas was directly blown and cooled. When the temperature change of the sample surface at that time was measured, the cooling rate at 800 ° C. to 600 ° C. was an average of 10 ° C./second. There was no change in the macro structure of the alloy, the grain size of the main phase, and the dispersion state of the eutectic region, but the minor axis of the phase with less R content precipitated in the eutectic region was about 3 μm. Next, a sintered magnet was produced from the obtained sample by the same method as in Example 1, and the maximum energy product was 37.5 MGOe and the coercive force was 18.0 kOe.
[0031]
(Example 6)
An alloy foil piece was prepared by the same method as in Example 1 so as to have the same composition as in Example 1, and a single roll type strip casting method using a copper roll having a diameter of 50 cm. In addition, the temperature measurement at the time of solidification is not implemented here. The roll peripheral speed at this time was 3 m / s. The thickness of the obtained alloy foil piece is as thin as 0.1 to 0.2 mm. The cross-sectional macrostructure of the roll contact surface side was a chill crystal, and the others were columnar crystals. After further polishing, the structure was observed by the same method as in Example 1. As a result, the main phase Nd of the columnar crystal part was observed. ₂ Fe ₁₄ The average particle size of the B phase was about 30 to 50 μm in the major axis direction and about 2 to 8 μm in the minor axis direction. The volume ratio of the columnar crystal portion to the entire ingot was 70%. The eutectic region is the main phase Nd. ₂ Fe ₁₄ The minor axis of the phase which is present mainly in the grain boundary portion of the B phase and has a smaller content of R precipitated in the eutectic region was 1 μm or less. Next, from the obtained ingot, a sintered magnet was produced in the same manner as in Example 1. The maximum energy product was 37.1 MGOe and the coercive force was 20.1 kOe.
[0032]
(Comparative Example 3)
The ingot produced in Comparative Example 2 was heat-treated at 900 ° C. for 1 hour in an argon atmosphere with a thickness of 20 mm, and then cooled by blowing argon gas directly. When the temperature change of the sample at that time was measured, the cooling rate at 800 ° C. to 600 ° C. was an average of 3 ° C./second. There was no change in the macro structure of the ingot, the particle size of the main phase, and the dispersion state of the eutectic region, and the particle size of the phase with less R content precipitated in the eutectic region was also 5 μm or more as before the heat treatment. It was. Next, a sintered magnet was produced from the obtained sample by the same method as in Example 1, and the maximum energy product was 34.9 MGOe and the coercive force was 15.1 kOe.
[0033]
(Comparative Example 4)
The ingot produced in Comparative Example 2 was heat-treated at 600 ° C. for 1 hour in an argon atmosphere with a thickness of 20 mm, and then cooled by blowing argon gas directly. There was no change in the macro structure of the ingot, the particle size of the main phase, and the dispersion state of the eutectic region, and the particle size of the phase with less R content precipitated in the eutectic region was also 5 μm or more as before the heat treatment. It was. Next, a sintered magnet was produced from the obtained sample in the same manner as in Example 1. The maximum energy product was 35.1 MGOe and the coercive force was 15.3 kOe.
[0034]
[Table 1]

[0035]
【The invention's effect】
INDUSTRIAL APPLICABILITY According to the present invention, it is possible to produce a raw material alloy that is optimal as a raw material for high-performance rare earth magnets without using complicated processes and equipment, which is extremely useful.
[Brief description of the drawings]
FIG. 1 is a diagram schematically showing the microstructure of an alloy according to the present invention.
FIG. 2 is a diagram schematically showing a microstructure of a conventional alloy.
FIG. 3 is a schematic diagram showing an enlarged structure of a eutectic region of an alloy according to the present invention.
FIG. 4 is a schematic diagram showing an enlarged structure of a eutectic region of a conventional alloy.
[Explanation of symbols]
1 R ₂ Fe ₁₄ B columnar crystals
2 Eutectic region
Phase with more 3R content
4 R phase with less content

Claims

Based on R ₂ Fe ₁₄ B columnar crystals and R ₂ Fe ₁₄ B crystals, which are composed of R (at least one of rare earth elements including Y), T (transition metal which essentially contains Fe), and B as main components. Also has a eutectic region with a high R content, and the eutectic region is composed of a portion with a relatively large amount of R and a portion with a smaller content of R crystallized in a rod shape. The phase of the minor axis direction of the phase with a lower content of 3 μm or less, the size of the major axis R ₂ Fe ₁₄ B columnar crystal in the major axis direction is 50 μm or more, and the dimension in the minor axis direction is 10 μm. A rare earth magnet alloy characterized in that the volume fraction of the crystal crystallization region is 70% or more.