JP2004214390A

JP2004214390A - Method for manufacturing rare-earth magnet, and raw alloy and powder for rare-earth magnet

Info

Publication number: JP2004214390A
Application number: JP2002381779A
Authority: JP
Inventors: Kenji Konishi; 謙治小西; Hiroshi Abe; 浩史阿部; Shinya Tabata; 進也田畑
Original assignee: Santoku Corp
Current assignee: Santoku Corp
Priority date: 2002-12-27
Filing date: 2002-12-27
Publication date: 2004-07-29
Anticipated expiration: 2022-12-27
Also published as: JP4133315B2

Abstract

<P>PROBLEM TO BE SOLVED: To provide a method for manufacturing rare-earth magnet by which the rare earth magnet of high magnetic characteristic is supplied easily and stably and an end material amount generated in the case of working a magnet is reduced, by improving the oxidation and yield of R component so as to expand a sintering temperature area in a method using a single alloy, and to provide alloy and powder useful for the method. <P>SOLUTION: The method for manufacturing comprises a process for preparing a material alloy (A) and a raw material (B). The material alloy (A) has the specific composition of R-B-M. In the material alloy (A), an R<SB>2</SB>Fe<SB>14</SB>B phase is a main phase and is 77 vol.%, an R-rich phase is 15 to 23 vol.%, 90 vol.% of the R<SB>2</SB>Fe<SB>14</SB>B phase consists of crystal grains of a specific grain diameter, and the axial lattice constant of the R<SB>2</SB>Fe<SB>14</SB>B phase is larger than that in a balanced state by 0.3 to 1%. The material alloy (B) has the specific composition of R-B-M. In the material alloy (B), a main phase R<SB>2</SB>Fe<SB>14</SB>B is 85 vol.%, and an R-rich phase is not larger than 15 vol.%. The method also comprises a process of crushing the alloys (A) and (B), a process of molding crushed powder, a process of sintering the molded matter, and a heat treatment process for treating a sintered matter with heat. <P>COPYRIGHT: (C)2004,JPO&NCIPI

Description

【０００１】
【発明の属する技術分野】
本発明は、希土類元素、ボロン及び鉄を含む優れた磁石特性を有する希土類磁石を、焼結熱処理条件の緩やかな管理により、容易に、且つ安定的に得ることができる希土類磁石の製造法、該製造法に主に利用可能な希土類磁石用原料合金及び粉末に関する。
【０００２】
【従来の技術】
近年、電子機器の小型高性能化に伴って、希土類磁石の需要は伸び続けている。特に、高い磁気特性をもち比較的安価なＲ−Ｆｅ−Ｂ系希土類磁石の生産量は増加し続けており、用途の拡大に伴って、更なる高特性化と厳密な特性の管理が要求されてきている。Ｒ−Ｆｅ−Ｂ系希土類磁石の内部組織には、主相である強磁性のＲ_２Ｆｅ_１４Ｂ相と、比較的融点が低く希土類元素を多く含む非磁性のＲ−ｒｉｃｈ相とが存在し、特に高特性磁石には、該Ｒ−ｒｉｃｈ相が微細に分散されていることが必要である。
このような組織を有する希土類磁石の製造法は大きく２つある。１つはＲ_２Ｆｅ_１４Ｂ相とＲ−ｒｉｃｈ相とを別々の合金から供給する方法であり、一般に２合金法と呼ばれている。２合金法では、焼結時に液相となって磁石の緻密化を促進させるＲ−ｒｉｃｈ相を独立して調整できるため、Ｒ−ｒｉｃｈ相が微細に分散した組織を有する合金を製造するための焼結可能な温度範囲を広くとることができる（例えば、特許文献１参照）。
もう１つの方法は、ストリップキャスティング法で鋳造した単一の合金を用いる方法である。ストリップキャスティング法では、合金の冷却速度が速いため組織全体が微細化され、合金内のＲ−ｒｉｃｈ相が微細に分散される。そのため、粉砕、焼結後のＲ−ｒｉｃｈ相の分散性も良好となり、磁石の高特性化が実現できる（例えば、特許文献２参照）。
【０００３】
２合金法の場合、Ｒ−ｒｉｃｈ相を供給する合金は、Ｒの含有量が通常４０〜６０質量％程度と高いため、酸素に対して極めて活性で、微粉砕時に合金が酸化し、特に残留磁束密度等の磁気特性の低下を招く。場合によっては微粉砕以降の磁石作製工程で発火を引き起こすことさえあり、作業員の危険性、製品歩留の低下が問題となっている。酸化、発火を防ぐ為、高価な設備により各工程での雰囲気の厳密な制御を試みているが十分とは言えない。また、Ｒ−ｒｉｃｈ相を供給する合金は、微粉砕時、Ｒの含有量が高く、強度の低い部分が選択的に、且つ異常に微粉化されて粉砕室の系外へ排出されるために歩留が低下する。更には回収される微粉末の組成が、仕込みの組成からずれて目的とする磁力特性が得られ難い。このため、酸化及び超微粉化による組成ずれを見越してＲの含有量を多めに仕込むことが行われている。しかし、このようなＲの含有量を多めに仕込む操作は、組成ずれを解消できても、前記酸化の問題や、歩留低下の問題解決にはならない。
前記ストリップキャスティング法で鋳造した単一の合金を用いる場合には、Ｒ−ｒｉｃｈ相が微細に分散した組織を有する磁石を製造するための焼結可能な最適な温度域が高くて狭い為、焼結、時効処理時の温度管理に非常に厳しい精度が要求される。しかし、実際には、焼結、時効処理時の炉内には温度分布があるため、合金温度の数℃のばらつきは避け難く、そのため得られる希土類磁石の磁気特性や、縮率にばらつきが生じる。縮率のばらつきが大きいと成形歩留が悪くなることから、希土類磁石の加工時に発生する端材量が増加する。
【０００４】
【特許文献１】
特公平６−２１３２４号公報
【特許文献２】
特許第２６３９６０９号公報
【０００５】
【発明が解決しようとする課題】
本発明の目的は、２合金法におけるＲ分の酸化、Ｒ分の歩留の悪さを改善し、また、従来の単一の合金を用いる方法における焼結温度域を広げ、高磁気特性の希土類磁石を、容易に、且つ安定して供給することができ、希土類磁石の加工時に発生する端材量を減少させることが可能な希土類磁石の製造法を提供することにある。
本発明の別の目的は、高磁気特性の希土類磁石を安定して供給するための製造法に有用である希土類磁石用原料合金及びその粉末を提供することにある。
【０００６】
【課題を解決するための手段】
本発明によれば、イットリウムを含む希土類金属元素から選ばれる少なくとも１種からなるＲ２７．６〜３５．０質量％と、ボロン０．９４〜１．３０質量％と、鉄を含む残部Ｍとからなる組成を有し、Ｒ_２Ｆｅ_１４Ｂ相を主相とし、Ｒ_２Ｆｅ_１４Ｂ相の含有割合が７７体積％以上、Ｒ−ｒｉｃｈ相の含有割合が１５〜２３体積％であり、Ｒ_２Ｆｅ_１４Ｂ相の９０体積％以上が、短軸方向の粒径が０．１〜５０μｍ、長軸方向の粒径が３０〜５００μｍの結晶粒からなり、Ｒ_２Ｆｅ_１４Ｂ相のｃ軸格子定数が、平衡状態より０．３〜１％大きい原料合金（Ａ）及び、イットリウムを含む希土類金属元素から選ばれる少なくとも１種からなるＲ２７．６〜３５．０質量％と、ボロン０．９４〜１．３０質量％と、鉄を含む残部Ｍとからなる組成を有し、主相としてのＲ_２Ｆｅ_１４Ｂ相の含有割合が８５体積％以上、Ｒ−ｒｉｃｈ相の含有割合が１５体積％未満である原料合金（Ｂ）とを準備する準備工程と、原料合金（Ａ）及び（Ｂ）を粉砕する粉砕工程と、得られた粉砕粉末を成型する成型工程と、得られた成型物を焼結する焼結工程と、得られた焼結物を時効処理する熱処理工程とを含むことを特徴とする希土類磁石の製造法が提供される。
また本発明によれば、前記希土類磁石用原料合金（Ａ）が提供される。
更に本発明によれば、前記原料合金（Ａ）を水素吸蔵崩壊して得た粉末であって、粒径３５５〜８５０μｍの粉末を５０質量％以上含み、重量平均粒径が３００〜８００μｍであることを特徴とする希土類磁石用原料合金粉末が提供される。
【０００７】
【発明の実施の形態】
以下、本発明を更に詳細に説明する。
本発明の希土類磁石用原料合金（Ａ）は、後述する本発明の製造法に主に有用な合金であるが、希土類磁石用であれば、その用途は必ずしも本発明の製造法のみに限定されない。該原料合金（Ａ）の組成は、イットリウムを含む希土類金属元素から選ばれる少なくとも１種からなるＲ２７．６〜３５．０質量％と、ボロン０．９４〜１．３０質量％と、鉄を含む残部Ｍとからなる組成を有する。
【０００８】
Ｒの希土類金属元素は特に限定されないが、ランタン、セリウム、プラセオジム、ネオジム、イットリウム、ジスプロシウム等が好ましく挙げられる。Ｒの含有割合が２７．６質量％未満では、α−Ｆｅ相が多く析出して磁気特性が低下し、３５．０質量％を超えると、焼結体内部のＲ−ｒｉｃｈ相の割合が高くなって耐食性が低下する。
ボロンの含有割合が０．９４質量％未満では、Ｒ_２Ｆｅ_１７相が析出してＲ_２Ｆｅ_１４Ｂ相の割合が減少するため残留磁束密度が低下し、１．３０質量％を超えると、Ｂ−ｒｉｃｈ相の割合が増加して磁気特性及び耐食性がともに低下する。
残部Ｍ中の鉄の含有割合は、通常５０質量％以上、好ましくは６０質量％以上である。残部Ｍは、鉄以外に、コバルト、アルミニウム、クロム、マンガン、マグネシウム、銅、錫、タングステン、ニオブ、ガリウム等の遷移金属や、炭素、珪素又はこれらの２種以上等を含むことができる。更に、原料合金（Ａ）には、その他、酸素、窒素等の工業生産上の不可避不純分が含まれていても良い。
【０００９】
本発明の原料合金（Ａ）の組織は、Ｒ_２Ｆｅ_１４Ｂ相を主相とし、Ｒ_２Ｆｅ_１４Ｂ相の９０体積％以上が、短軸方向の粒径が０．１〜５０μｍ、長軸方向の粒径が３０〜５００μｍの結晶粒からなる。原料合金（Ａ）は、Ｒ_２Ｆｅ_１４Ｂ相を７７体積％以上、好ましくは７７〜８５体積％含有し、Ｒ−ｒｉｃｈ相を１５〜２３体積％、好ましくは１８〜２３体積％含有する。該Ｒ−ｒｉｃｈ相は、前記Ｒ_２Ｆｅ_１４Ｂ相を主相とする結晶粒を取り囲むように微細に分散されていることが好ましい。Ｒ−ｒｉｃｈ相の一部はＲ_２Ｆｅ_１４Ｂ相結晶粒内に微細に分散されて存在する。Ｒ−ｒｉｃｈ相の含有割合が１５体積％未満では、後述する本発明の製法に用いる場合、焼結温度域を広げる効果が小さく、２３体積％を超えると残留磁束密度が低下する。
原料合金（Ａ）における前記結晶粒の短軸方向及び長軸方向の粒径の測定は、薄片等の原料合金（Ａ）の断面組織を、光学顕微鏡により撮影した写真から測定することができる。また、前記各結晶相の体積率はＥＰＭＡの画像解析により求めることができる。
本願発明においては、日本電子（株）社製ＪＸＡ８８００を用いて合金鋳片の厚み方向中央部断面のＣｏｍｐ像を撮影し、画像処理ソフト（全自動粒子解析プログラムＸＭ−８７５６２）により、Ｃｏｍｐ像を解析してそれぞれの相の面積率を求め、体積率とした。ここで得られる値は、合金断面を２次元で解析した値である為、実際の体積率とは異なる場合がある。
【００１０】
原料合金（Ａ）のＲ_２Ｆｅ_１４Ｂ相のｃ軸格子定数は、後述する平衡状態より０．３〜１％大きく、好ましくは０．３〜０．８％大きい。ｃ軸格子定数は、合金を２５μｍ以下に粉砕して、粉末Ｘ線回折法により求めた。原料合金（Ａ）が、このような物性を示すのは、Ｒ_２Ｆｅ_１４Ｂ相の結晶格子中にＲ原子が侵入し、結晶格子が歪んでいる状態であるからだと考えられる。
前記ｃ軸格子定数の増大が平衡状態のｃ軸格子定数と比較して０．３％未満の場合、Ｒ−ｒｉｃｈ相の体積率が小さく、Ｒ_２Ｆｅ_１４Ｂ相格子中に侵入するＲ原子が少ないため、焼結時にＲ−ｒｉｃｈ相及びＲ_２Ｆｅ_１４Ｂ相より排出されるＲ分が少なく、後述する液相が十分供給されない。その為、焼結性が低下し、保持力が低下する。また、後述する本発明の製造法に用いた場合、焼結温度域を広げる効果が小さく、得られる希土類磁石の磁気特性や、縮率のばらつきを小さくできない。一方、ｃ軸格子定数が平衡状態のｃ軸格子定数と比較して１％を超える合金の製造は困難である。
【００１１】
前記平衡状態の合金とは、合金を６００〜８００℃で４０時間以上熱処理し、その後、１℃／分以下の冷却速度で徐冷して得た合金をいう。
【００１２】
原料合金（Ａ）は、従来のストリップキャスト法により鋳造した原料合金と比較してＲ−ｒｉｃｈ相の体積率は高く、且つ該Ｒ−ｒｉｃｈ相が非常に微細に分散した状態で存在する。該Ｒ−ｒｉｃｈ相は、従来のストリップキャスト法により鋳造した原料合金のＲ−ｒｉｃｈ相と比較して、Ｆｅ、Ｂが多く存在する状態であると考えられる。このような組成及び結晶組織を有する原料合金（Ａ）は、後述する本発明の製造法に用いることにより、焼結の際、Ｒ−ｒｉｃｈ相及びＲ_２Ｆｅ_１４Ｂ相の結晶格子中に侵入しているＲ原子が、従来の２合金法においてＲ−ｒｉｃｈ相を供給する粒界相合金のように液相を作り出す。Ｒが微細に分散したＲ−ｒｉｃｈ相、並びにＲ_２Ｆｅ_１４Ｂ相から均一に排出される為、その液相は、従来の２合金法の粒界相合金と比較して、焼結に必要な箇所に必要量のみがより均一に供給されることとなる。また、原料合金（Ａ）は、２合金法における粒界相合金と較べてＲ量が少ない為、非磁性相であるＲ−ｒｉｃｈ相を必要以上に生成させず、残留磁束密度を低下させない。
【００１３】
原料合金（Ａ）は、その形態が、板厚０．０３〜２．００ｍｍ、好ましくは０．１０〜１．０ｍｍの薄片であることが望ましい。
原料合金（Ａ）は、合金単位質量当りのＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量が、平衡状態のＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量より３〜２２％大きく、好ましくは１２〜２２％大きい。また、Ｒ−ｒｉｃｈ相の水素吸蔵量が、平衡状態のＲ−ｒｉｃｈ相の水素吸蔵量より１０〜６０％小さく、好ましくは１０〜３０％小さい。以下、平衡状態より水素吸蔵量が大きい場合は＋、小さい場合は−を付して記載する。
このような水素吸蔵特性は、原料合金（Ａ）のＲ_２Ｆｅ_１４Ｂ相及びＲ−ｒｉｃｈ相の組成の特徴に起因するものと考えられる。原料合金（Ａ）のＲ_２Ｆｅ_１４Ｂ相は、Ｒ成分を通常より多く含み、水素吸蔵性がよく、Ｒ−ｒｉｃｈ相はＲ以外の成分（Ｆｅ、Ｂ等）を通常より多く含み、これらの成分はＲと比較的水素吸蔵性が低い化合物を形成し、その結果、Ｒ−ｒｉｃｈ相の体積率は高いものの水素吸蔵量が比較的低いものとなっていると考えられる。
Ｒ_２Ｆｅ_１４Ｂ相の水素吸蔵量が、平衡状態のＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量＋３．０％より小さく、またＲ−ｒｉｃｈ相の水素吸蔵量が、平衡状態のＲ−ｒｉｃｈ相の水素吸蔵量と比較して−１０％より大きい場合は、Ｒ_２Ｆｅ_１４Ｂ相が粗大となり、焼結性が低下して保磁力が低下する恐れがあるので好ましくない。Ｒ_２Ｆｅ_１４Ｂ相の水素吸蔵量が、平衡状態のＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量と比較して＋２２％より大きく、またＲ−ｒｉｃｈ相の水素吸蔵量が、平衡状態のＲ−ｒｉｃｈ相の水素吸蔵量と比較して−６０％より小さい場合は、合金の製造が困難である。
ここで、平衡状態とは前述のとおりであり、水素吸蔵量は、ＰＣＴ装置により以下のようにして測定することができる。
【００１４】
まず、一定容量のセル内に測定する合金を投入し、１〜３Ｐａに真空引きを行う。減圧雰囲気で４０℃の温度に保持し、４気圧の水素雰囲気にする。次に、合金の水素吸蔵が停止した時点での水素圧を測定する。この時、水素吸蔵停止時の水素圧力が１気圧以下である場合、再び４気圧の水素雰囲気に置く。水素吸蔵停止時の圧力が１気圧以上になるまでこれを繰り返す。これらの水素圧の変動の合計から合金の水素吸蔵量を求め、セル内に投入した合金重量から、単位重量当たりの水素吸蔵量を求めることができる。
次に、同じ鋳造条件により得られたＲ添加量の異なる合金サンプルについても上記測定を行い、横軸がＲ添加量、縦軸が水素吸蔵量のグラフを作成する。そのグラフの近似曲線からＲ_２Ｆｅ_１４Ｂ相の量論組成のＲ量（２６．７質量％）での水素吸蔵量を求め、それをＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量とする。各組成での合金全体の水素吸蔵量とＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量との差よりＲ−ｒｉｃｈ相の水素吸蔵量を求めることができる。
平衡状態の合金のＲ_２Ｆｅ_１４Ｂ相の水素吸蔵量及びＲ−ｒｉｃｈ相の水素吸蔵量も同様な方法により求めることができる。
【００１５】
原料合金（Ａ）は、Ｒ−ｒｉｃｈ相の体積率（Ｒｘ）を、Ｒｘ＝Ｒｃ×Ｘ（ここで、Ｒｃは平衡状態でのＲ−ｒｉｃｈ相の体積率である）と表した際のＸが２．２〜５．０、特に３．０〜５．０であることが好ましい。Ｒ−ｒｉｃｈ相の体積率（Ｒｘ）及び平衡状態でのＲ−ｒｉｃｈ相の体積率（Ｒｃ）は、ＥＰＭＡの画像解析により求めることができる。
Ｘが５．０を超える場合、Ｒ−ｒｉｃｈ相の体積割合が大きくなり、得られる磁石の残留磁束密度が低下する場合があり、Ｘが２．２未満では、前述に液相が十分供給されず、焼結性が低下し保磁力が低下する場合があるので好ましくない。
【００１６】
本発明の希土類磁石用原料合金粉末は、前述の本発明の原料合金（Ａ）を水素吸蔵崩壊して得た粉末であって、後述する本発明の製造法に有用な合金粉末であるが、希土類磁石用であれば、その用途は必ずしも本発明の製造法のみに限定されない。
本発明の希土類磁石用原料合金粉末は、粒径３５５〜８５０μｍの粉末を５０質量％以上含み、重量平均粒径が３００〜８００μｍ、好ましくは３００〜５００μｍの粉末である。
前記水素吸蔵崩壊は、例えば、原料合金（Ａ）を、水素圧力０．１〜０．４ＭＰａ、常温〜１００℃の雰囲気中で３０分以上水素吸蔵させた後、３００〜５００℃で１Ｐａ以下の雰囲気になるまで真空引きを行って、脱水素させる方法等により行うことができる。
水素吸蔵崩壊させて得た本発明の希土類磁石用原料合金粉末は、極端な微粉が少なく、酸化が抑制され、Ｒ分の歩留が良く、また極端に粗大な粉砕粉が少なく、後の微粉砕工程が効率良く行うことができる。
【００１７】
本発明の原料合金（Ａ）を調製するには、例えば、以下の方法等により得ることができる。
まず、上述の原料合金（Ａ）の組成に調整した、Ｒ、ボロン及びＭの原料金属や母合金を、不活性ガス雰囲気下、高周波溶融法により溶融した後、該溶融物を単ロール、双ロール又はディスク等を用いるストリップキャスティング法により連続的に凝固させる。その際、合金の融点からロール剥離までの平均冷却速度は、５０〜３０００℃／秒にて行うことができる（１次冷却工程）。該１次冷却工程により、Ｒ_２Ｆｅ_１４Ｂ相の短軸方向及び長軸方向の粒径、Ｒ_２Ｆｅ_１４Ｂ相及びＲ−ｒｉｃｈ相の体積割合、Ｒ_２Ｆｅ_１４Ｂ相のｃ軸格子定数のおおよそが確定する。
平均冷却速度が速すぎるとチル晶が発生し、規定の短軸方向、長軸方向の粒径のＲ_２Ｆｅ_１４Ｂ相を主相とする結晶粒を規定の体積率とすることができず、残留磁束密度が低下する。また遅すぎるとＲ_２Ｆｅ_１４Ｂ相のｃ軸格子定数が規定の範囲とならず、Ｒ_２Ｆｅ_１４Ｂ相が粗大化し、Ｒ−ｒｉｃｈ相の分散性が低下し、保磁力が低下する。
【００１８】
更に、ロール剥離後の鋳片の冷却過程を制御することにより、Ｒ_２Ｆｅ_１４Ｂ相のｃ軸格子定数及び水素吸蔵量を好ましく制御した合金を均一に製造することができる。例えば、前記１次冷却工程後、鋳片を３０℃／秒以上の平均冷却速度で三元共晶点＋３０℃の温度まで冷却し（２次冷却工程）、三元共晶温度±３０℃の範囲を３０秒以内で冷却させる（３次冷却工程）方法が挙げられる。該３次冷却工程は、好ましくは５秒以内で行う。該工程は、Ｒ−ｒｉｃｈ相内の析出相を制御して水素吸蔵特性を制御するために好ましく実施できる。３次冷却工程において、三元共晶温度±３０℃の範囲における保持時間が長すぎると、Ｒ_２Ｆｅ_１４Ｂ相が粗大化し、Ｒ−ｒｉｃｈ相の偏析が生じ保磁力の低下を招く恐れがある。その後、好ましくは５〜３０℃／分で１００℃以下まで冷却する（４次冷却工程）ことにより本発明の原料合金（Ａ）を得ることができる。
前記原料合金（Ａ）の製造における温度制御は、ロール離脱後の鋳片が接触する鋳造装置の部材の選定、加熱機構又は冷却機構を有する温度制御装置等により適宜行うことができる。
【００１９】
本発明の希土類磁石の製造法は、前記原料合金（Ａ）及び、イットリウムを含む希土類金属元素から選ばれる少なくとも１種からなるＲ２７．６〜３５．０質量％と、ボロン０．９４〜１．３０質量％と、鉄を含む残部Ｍとからなる組成を有し、主相としてのＲ_２Ｆｅ_１４Ｂ相の含有割合が８５体積％以上、Ｒ−ｒｉｃｈ相の含有割合が１５体積％以下である原料合金（Ｂ）とを準備する準備工程と、原料合金（Ａ）及び（Ｂ）を粉砕する粉砕工程と、得られた粉砕粉末を成型する成型工程と、得られた成型物を焼結する焼結工程と、得られた焼結物を時効処理する熱処理工程とを含む。
【００２０】
準備工程に用いる原料合金（Ａ）は、前述の原料合金（Ａ）を用いることができ、また、原料合金（Ａ）として、前述の本発明の希土類磁石用原料合金粉末を用いることもできる。
準備工程に用いる原料合金（Ｂ）の組成範囲は、原料合金（Ａ）と同じであり、必要により含有していても良い元素等も上述した原料合金（Ａ）と同じものが挙げられる。従って、原料合金（Ａ）及び（Ｂ）の組成は、規定する組成範囲において同一でも異なっていても良い。
【００２１】
原料合金（Ｂ）のＲ_２Ｆｅ_１４Ｂ相の体積率は８５％以上であれば良く、該体積率は上述と同様な方法で求めることができる。
原料合金（Ｂ）の結晶組織は、Ｒ−ｒｉｃｈ相の含有割合が原料合金（Ａ）と異なる以外は同一であっても異なっていても良い。原料合金（Ｂ）におけるＲ−ｒｉｃｈ相の含有割合は、１５体積％以下、好ましくは３〜１０体積％である。原料合金（Ｂ）のＲ−ｒｉｃｈ相の含有割合が１５体積％を超える場合は、残留磁束密度が低下する。
【００２２】
準備工程において準備する原料合金（Ａ）は、上述の通り、液相を最適な条件で供給し、同時にＲ_２Ｆｅ_１４Ｂ相も供給する。一方、原料合金（Ｂ）は、従来の２合金法における主相合金と同様に、主にＲ_２Ｆｅ_１４Ｂ相を供給する役割を担う。原料合金（Ａ）及び（Ｂ）は、従来の２合金法のようにＲ含有量の高い合金を使用しない為、Ｒ分の酸化を抑制し、Ｒ分の歩留が良い。原料合金（Ａ）の作用により、単一合金を用いる場合に比較し、最適焼結温度幅を広げることができる為、得られる希土類磁石の磁気特性及び縮率のばらつきを抑制できる。
【００２３】
準備工程において、原料合金（Ａ）及び（Ｂ）の混合割合は、求める希土類磁石の磁気特性、原料合金（Ａ）の特徴、原料合金（Ｂ）の特徴により適宜選択することができるが、質量比で原料合金（Ａ）：（Ｂ）＝１：１〜３０が好ましい。原料合金（Ｂ）は、規定の範囲内で適当な組成、若しくは結晶組織を有するものが使用できる。例えば、高残留磁束密度の永久磁石を得る場合には、原料合金（Ｂ）としてＲ含有割合が２８．１〜３０．０質量％の合金を用いることが好ましい。このようなＲ含有割合であれば、主相の体積率が大きくなり残留磁束密度が向上する。またＲ含有割合が３０．０〜３３．０質量％の合金を使用する場合、Ｒ_２Ｆｅ_１４Ｂ相の短軸方向の平均粒径が５μｍ以上であれば、製造工程において粉砕した場合、粉砕粉中に結晶方位の異なる２つ以上の主相を含む割合が減少する為、高残留磁束密度の磁気特性が得られる。
また高保磁力の永久磁石を得る場合には、原料合金（Ｂ）としてＲ含有割合が３１〜３５質量％の合金の使用が好ましい。このようなＲ含有割合であれば、Ｒ含有割合が大きいほどＲ−ｒｉｃｈ相の体積率が増加して焼結性が向上し、保磁力が向上する。また原料合金（Ｂ）のＲ含有割合が３１質量％未満であっても、Ｒ_２Ｆｅ_１４Ｂ相の単軸方向の平均粒径が４．０μｍ以下で、Ｒ−ｒｉｃｈ相の含有割合が１５体積％程度であれば、焼結性が向上して高保磁力が得られる。このような原料合金（Ｂ）を用いる場合の原料合金（Ａ）の混合割合は、全原料合金に対して５０質量％以下が好ましい。
【００２４】
準備工程において、原料合金（Ａ）は、上述の方法等により得ることができ、一方、原料合金（Ｂ）も製造条件を適宜変えることにより得ることができる。例えば、組成を変更したり、冷却工程において昇温、保持工程を含む冷却速度を制御すること等により得ることができる。
原料合金（Ａ）及び（Ｂ）は、ストリップキャスト法で得た薄片の形状、薄片を粗粉砕した粗粉砕粉の形状であっても良く、このような粗粉砕形状のものとしては、例えば、本発明の希土類磁石用原料合金粉末が挙げられる。また、前記規定の範囲の原料合金（Ｂ）であれば、金型鋳造法で得たインゴットを粗粉砕した粗粉砕粉の形状のものであっても良い。
【００２５】
本発明の製造法では、次に、粉砕工程を行う。該粉砕工程では、通常、水素化粉砕を行った後、ジェットミル等の粉砕機を用いて、原料合金（Ａ）及び（Ｂ）を平均粒径３〜６μｍ程度に粉砕する。該粉砕は、原料合金（Ａ）及び（Ｂ）を混合した状態で行うことが作業上、もしくは酸化を抑制する面で好ましい。しかし、それぞれを粉砕後、粉砕粉同士を混合することも可能である。
【００２６】
本発明の製造法では、次いで、得られた粉砕粉末を所望の形状及び大きさに成型する。成型は、希土類磁石製造に採用される公知の方法で行うことができ、例えば、磁場中において加圧して成型する方法等により行うことができる。通常は、１５〜３０ｋＯｅの磁場中にて０．５〜３．０ｔ／ｃｍ^２の圧力で成型する。
【００２７】
本発明の製造法では、前記成型物を焼結するための焼結工程を行う。該焼結工程における焼結温度は、通常１０００〜１１００℃の範囲から適宜選択することができ、焼結時間も通常１〜５時間の範囲から適宜選択することができるが、本発明の製造法における焼結温度域は、上述の原料合金（Ａ）を採用するので、従来の単一の合金を用いた場合と比較すると広い範囲で設定しても得られる希土類磁石の磁気特性及び縮率のばらつきを小さくすることができる。従って、焼結温度管理が従来の単一の合金を用いたより緩和される。
【００２８】
本発明の製造法では、前記焼結物を時効処理する熱処理工程を行う。該時効処理も所望の希土類磁石を得るために公知の方法から適宜条件を選択して行うことができる。該時効処理は、例えば、４５０〜９５０℃の温度範囲から２回以上に分けて温度を下げて所望時間保持する方法等により行うことができる。
【００２９】
【実施例】
以下、実施例及び比較例により、本発明を更に詳細に説明するが、本発明はこれらに限定されない。
実施例１
（原料合金（Ａ）の調製）
Ｎｄ３２．０質量％、Ｄｙ１．０質量％、Ｂ１．００質量％、Ａｌ０．２０質量％、Ｃｏ１．０質量％、残部鉄になるように、ネオジムメタル、ジスプロシウムメタル、フェロボロン、アルミニウム、コバルト及び鉄を配合し、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た。次いで、得られた合金溶湯を、水冷式の銅製単ロール鋳造装置を用いてストリップキャスティング法により鋳造し、厚さ約０．２ｍｍの鋳片を得た。この合金の三元共晶点は約６４０℃である。この合金のロールに接する直前の溶湯の温度は約１３５０〜１４００℃で、ロールから剥離した直後の鋳片の温度を赤外線熱画像計測装置で測定したところ約６００℃であった。ロール上での冷却時間は約０．６秒であった。
次に、ロール剥離後の鋳片を回転ドラム式の水冷装置により冷却し、４０分後回収し、原料合金（Ａ）としての試料１を得た。ロール剥離後、水冷装置に入るまでに要した時間は約０．８秒であった。水冷装置に鋳片を移動した直後の鋳片温度は約４５０℃、取り出し直後の鋳片温度は約６０℃であった。
【００３０】
得られた鋳片（試料１）の断面組織を光学顕微鏡により撮影し、Ｒ_２Ｆｅ_１４Ｂ相の短軸方向の平均粒径、長軸方向の平均粒径を測定し、それぞれの平均粒径を求め、ＥＰＭＡの画像解析によりＲ_２Ｆｅ_１４Ｂ相の体積率及びＲ−ｒｉｃｈ相の体積率を求めた。短軸方向の平均粒径が０．１〜５０μｍで、且つ長軸方向の平均粒径が３０〜５００μｍである結晶粒（Ｘ）の体積率を求めた。短軸方向の粒径は３．３μｍ、長軸方向の粒径は７４μｍ、Ｒ_２Ｆｅ_１４Ｂ相の体積率は８２体積％、結晶粒（Ｘ）の体積率は９５体積％、Ｒ−ｒｉｃｈ相の体積率は１８体積％であった。また、鋳片を約２５μｍ程度に粉砕後、Ｘ線回折装置にてＲ_２Ｆｅ_１４Ｂ相のｃ軸格子定数を測定したところ１２．３４Åであった。
得られた鋳片を３０℃、０．１ＭＰａの水素雰囲気中で１時間水素化した後、４００℃で脱水素することで水素粉砕を行い、得られた粉砕粉末を、ロータップ式標準篩振盪機で篩い分けしたところ、粒径３５５〜８５０μｍの粉末が約７４質量％、粉末の平均粒径は約４５０μｍであった。更に、ＰＣＴ装置により水素吸蔵量を求めたところ０．３９３質量％であった。該吸蔵量から主相のＲ_２Ｆｅ_１４Ｂ相及びＲ−ｒｉｃｈ相の水素吸蔵量を求めたところ、主相へ約０．２７８質量％、Ｒ−ｒｉｃｈ相へ０．１１５質量％吸蔵されていた。
得られた鋳片を８００℃で４０時間熱処理を行い、平衡状態とし、１℃／分以下の冷却速度で冷し、上述の方法で分析を行ったところ、主相のＲ_２Ｆｅ_１４Ｂ相のｃ軸格子定数は１２．２５Å、水素吸蔵量は０．４０８質量％、主相の水素吸蔵量は０．２３７質量％、Ｒ−ｒｉｃｈ相の水素吸蔵量は０．１７１質量％、Ｒ−ｒｉｃｈ相の体積率は約４％であった。
以上の測定結果を表１及び２に示す。
【００３１】
（原料合金（Ｂ）の調整）
Ｎｄ３２．０質量％、Ｄｙ１．０質量％、Ｂ１．００質量％、Ａｌ０．２０質量％、Ｃｏ１．０質量％及び残部鉄になるように、ネオジムメタル、ジスプロシウムメタル、フェロボロン、アルミニウム、コバルト、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た。次いで、水冷式の銅製単ロール鋳造装置を用いてストリップキャスティング法により鋳造し、厚さ約０．４ｍｍの鋳片を得た。ロールに接する直前の溶湯の温度は１３００〜１３５０℃で、ロールから剥離した直後の鋳片の温度を赤外線熱画像計測装置で測定したところ約８００℃であった。ロール上での冷却時間は約１．２秒であった。
次に、ロール剥離後の鋳片を鋼鉄製の容器に回収し、容器を密閉後、大気中に取り出して放冷し、１５００分後に回収し、原料合金（Ｂ）としての試料１ａを得た。容器内の鋳片の温度は、回収直後で約６６５℃、回収後８０分経過後で約６１５℃、１５００分経過後に鋳片を取り出したときは約９０℃であった。またロール剥離後、鋼鉄製の容器に入るまでに要した時間は約０．８秒であった。試料１ａについて上述の試料１と同様の分析を行った。結果を表１及び２に示す。また、試料１及び試料１ａ製造時の熱履歴を表すグラフを図１に示す。
【００３２】
（永久磁石の製法）
上記で調製した試料１及び試料１ａを５：５の質量比でドラムミキサーに導入して混合した。３０℃、０．１ＭＰａの水素雰囲気中で１時間水素化した後、４００℃で脱水素することで水素粉砕を行った。次いで、ジェットミルにより平均粒径が５．０μｍになるように粉砕を行った。
次いで、１５ｋＯｅの磁場中にて２．５ｔｏｎ／ｃｍ^２の圧力で成型を行い、得られた成型体を真空中で４時間焼結した。その際、焼結温度を１０５５℃、１０６０℃及び１０６５℃と変えた。焼結後、１段目の熱処理を９００℃で１時間、２段目の熱処理を５００℃で２時間行い時効処理した。得られた永久磁石の磁気特性を常法により測定した。結果を表３に示す。
また、得られた永久磁石の配向収縮率を以下の定義に従い測定した。一般に、Ｒ−Ｆｅ−Ｂ系焼結磁石では、磁気的に異方化させるため、磁界中で粒子を配向させながらプレス成型を行う。これに伴い、焼結時の収縮量は粒子の配向方向（ｃ軸方向）とそれに垂直なａ軸方向とで異なる。配向方向の収縮量をΔＬ、焼結及び時効
処理前の成型体の長さをＬ_０とした場合、配向収縮率は以下の式により求めることができる。結果を表３に示す。
収縮率＝ΔＬ／Ｌ_０
【００３３】
実施例２〜４
実施例１で調製した試料１及び試料１ａの混合比を、表２に示すとおり代えた以外は実施例１と同様にして永久磁石を調製し、各磁気特性等を測定した。結果を表３に示す。
【００３４】
実施例５
合金組成がＮｄ３４．０質量％、Ｄｙ１．０質量％、Ｂ１．００質量％、Ａｌ０．２０質量％、Ｃｏ１．０質量％及び残部鉄になるようにした以外は、実施例１における試料１と同様な方法により、原料合金（Ａ）である試料２を調製した。この合金の三元共晶点は約６４０℃である。鋳造中の鋳片の熱履歴は試料１とほぼ同じであった。
また合金の組成がＮｄ３１．５質量％、Ｄｙ１．０質量％、Ｂ１．０質量％、Ａｌ０．２０質量％、Ｃｏ１．０質量％及び残部鉄になるようにした以外は、実施例１における試料１ａと同様な方法により、原料合金（Ｂ）である試料２ａを調製した。この合金の三元共晶点は約６４０℃である。鋳造中の鋳片の熱履歴は試料１ａとほぼ同じであった。得られた試料２及び試料２ａについて、実施例１における試料１と同様な分析を行った。結果を表１及び２に示す。
更に、試料２及び試料２ａを２：８の質量比で混合した以外は実施例１と同様な方法により永久磁石を調製し、磁気特性等を測定した。結果を表３に示す。
【００３５】
比較例１
実施例１で調製した試料１ａのみを用いた以外は実施例１と同様に永久磁石を調製し、磁気特性等を測定した。結果を表３に示す。
【００３６】
比較例２
合金組成が、Ｎｄ３０．５質量％、Ｂ１．１１質量％、Ａｌ０．２０質量％及び残部鉄になるようにした以外は、実施例１における試料１ａと同様に、原料合金（Ｂ）としての試料３ａを調製した。この合金の三元共晶点は約６４０℃である。鋳造中の鋳片の熱履歴は試料１ａとほぼ同じであった。
また合金組成が、Ｎｄ４５．５質量％、Ｄｙ１０．０質量％、Ａｌ０．２０質量％、Ｃｏ１０．０質量％及び残部鉄になるようにし、熱履歴を以下のようにした以外は、実施例１における試料１ａと同様な方法により、原料合金（Ａ）及び（Ｂ）とは異なる試料４ａを得た。この際、ロールに接する直前の溶湯の温度は約１４５０〜１５００℃で、ロールから剥離した直後の鋳片の温度は約５５０℃であった。ロール上での冷却時間は約１．８秒であった。更に、ロール剥離後の鋳片を回転ドラム式の水冷装置により冷却し、４０分経過後に回収した。水冷装置に鋳片を移動した直後の鋳片温度は約４００℃、取り出し直後の鋳片温度は約５５℃であった。試料３ａ及び試料４ａの鋳片厚みはいずれも約０．４ｍｍであった。
試料３ａについて、実施例１における試料１と同様な分析を行った。結果を表１及び２に示す。
また、得られた試料３ａ及び試料４ａを９：１の質量比で混合した以外は実施例１と同様に永久磁石を調製し、磁気特性等を測定した。結果を表３に示す。
【００３７】
【表１】

【００３８】
【表２】

【００３９】
【表３】

【００４０】
【発明の効果】
本発明の希土類磁石の製造法では、特に、本発明の原料合金（Ａ）や本発明の希土類磁石原料合金を、原料合金（Ｂ）と混合して用いるので、製造時における焼結温度管理が、従来の単一の合金を使用して希土類磁石を製造する方法より緩和され、更に、従来の２合金法におけるＲ成分の酸化問題や歩留の悪さが改善される。しかも、得られる希土類磁石は、焼結温度のばらつきによる磁気特性の低下、並びに縮率のばらつきがなく、従来の２合金法により調製した永久磁石より、磁気特性に優れる。
また本発明の希土類磁石用原料合金（Ａ）及びその粉末は、特定の組成及び組織を有するので、高磁気特性の希土類磁石を安定して容易に得るための製造法に有用であり、特に、本発明の前記製造法に有用である。
【図面の簡単な説明】
【図１】
実施例１で調製した試料１及び試料１ａの製造時の熱履歴を表すグラフである。[0001]
TECHNICAL FIELD OF THE INVENTION
The present invention provides a method for manufacturing a rare earth magnet which can easily and stably obtain a rare earth magnet having excellent magnet properties including a rare earth element, boron and iron, by moderate management of sintering heat treatment conditions. The present invention relates to raw material alloys and powders for rare earth magnets that can be mainly used in production methods.
[0002]
[Prior art]
In recent years, demand for rare-earth magnets has been increasing along with miniaturization and high performance of electronic devices. In particular, the production of relatively inexpensive R-Fe-B rare earth magnets having high magnetic properties continues to increase, and with the expansion of applications, further higher properties and strict control of properties are required. Is coming. The internal structure of the R—Fe—B based rare earth magnet includes a ferromagnetic R₂Fe₁₄There is a B phase and a nonmagnetic R-rich phase having a relatively low melting point and containing a large amount of rare earth elements. Particularly, high-performance magnets require that the R-rich phase be finely dispersed. .
There are roughly two methods for producing a rare earth magnet having such a structure. One is R₂Fe₁₄This is a method in which the B phase and the R-rich phase are supplied from different alloys, and is generally called a two-alloy method. In the two-alloy method, since the R-rich phase which becomes a liquid phase during sintering and promotes the densification of the magnet can be adjusted independently, it is necessary to produce an alloy having a structure in which the R-rich phase is finely dispersed. A sinterable temperature range can be widened (for example, see Patent Document 1).
Another method is to use a single alloy cast by a strip casting method. In the strip casting method, since the cooling rate of the alloy is high, the entire structure is refined, and the R-rich phase in the alloy is finely dispersed. Therefore, the dispersibility of the R-rich phase after pulverization and sintering is also improved, and high characteristics of the magnet can be realized (for example, see Patent Document 2).
[0003]
In the case of the two-alloy method, the alloy supplying the R-rich phase is extremely active against oxygen because the R content is usually as high as about 40 to 60% by mass. This causes a decrease in magnetic properties such as magnetic flux density. In some cases, ignition may be caused even in the magnet manufacturing process after the pulverization, which poses a problem of danger to workers and a reduction in product yield. In order to prevent oxidation and ignition, strict control of the atmosphere in each step is attempted with expensive equipment, but this is not sufficient. In addition, the alloy that supplies the R-rich phase has a high R content during pulverization, and a low-strength portion is selectively and abnormally pulverized and discharged to the outside of the pulverization chamber. Yield decreases. Further, the composition of the fine powder to be recovered deviates from the composition of the preparation, and it is difficult to obtain the desired magnetic force characteristics. For this reason, a large amount of R is charged in anticipation of a composition shift due to oxidation and ultrafine powdering. However, such an operation of charging a large amount of R does not solve the problem of oxidation and the problem of reduction in yield, even if the composition deviation can be eliminated.
When a single alloy cast by the strip casting method is used, the optimum temperature range in which sintering for producing a magnet having a structure in which the R-rich phase is finely dispersed is high and narrow, so As a result, very strict accuracy is required for temperature control during aging treatment. However, in practice, since there is a temperature distribution in the furnace during sintering and aging, variations in the alloy temperature of several degrees Celsius are unavoidable, and therefore the magnetic properties and shrinkage of the obtained rare earth magnets vary. . If the variation of the shrinkage ratio is large, the forming yield is deteriorated, so that the amount of offcuts generated during the processing of the rare earth magnet increases.
[0004]
[Patent Document 1]
Japanese Patent Publication No. Hei 6-21324
[Patent Document 2]
Japanese Patent No. 2639609
[0005]
[Problems to be solved by the invention]
An object of the present invention is to improve the oxidization of R component and the poor yield of R component in the two-alloy method, widen the sintering temperature range in the conventional method using a single alloy, and obtain a rare earth element having high magnetic properties. It is an object of the present invention to provide a method for manufacturing a rare earth magnet, which can easily and stably supply a magnet and can reduce the amount of offcuts generated during processing of the rare earth magnet.
Another object of the present invention is to provide a rare earth magnet raw material alloy and powder thereof which are useful for a production method for stably supplying a rare earth magnet having high magnetic properties.
[0006]
[Means for Solving the Problems]
According to the present invention, R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements including yttrium, 0.94 to 1.30% by mass of boron, and the balance M including iron Having a composition consisting of₂Fe₁₄B phase as main phase, R phase₂Fe₁₄The content ratio of the B phase is 77% by volume or more, the content ratio of the R-rich phase is 15 to 23% by volume,₂Fe₁₄90% by volume or more of the B phase is composed of crystal grains having a particle diameter in the short axis direction of 0.1 to 50 μm and a particle diameter in the long axis direction of 30 to 500 μm.₂Fe₁₄A raw material alloy (A) in which the c-axis lattice constant of the B phase is 0.3 to 1% larger than the equilibrium state, and R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements including yttrium. And a composition consisting of 0.94 to 1.30 mass% of boron and the balance M containing iron.₂Fe₁₄A preparation step of preparing a raw material alloy (B) having a B phase content of 85% by volume or more and an R-rich phase content of less than 15% by volume, and pulverizing the material alloys (A) and (B) A crushing step, a molding step of molding the obtained crushed powder, a sintering step of sintering the obtained molded article, and a heat treatment step of aging the obtained sintered article. A method for manufacturing a rare earth magnet is provided.
Further, according to the present invention, there is provided the raw alloy (A) for a rare earth magnet.
Further, according to the present invention, the powder obtained by subjecting the raw material alloy (A) to hydrogen absorption and collapse contains at least 50% by mass of a powder having a particle size of 355 to 850 μm and a weight average particle size of 300 to 800 μm. A raw material alloy powder for a rare earth magnet is provided.
[0007]
BEST MODE FOR CARRYING OUT THE INVENTION
Hereinafter, the present invention will be described in more detail.
The rare-earth magnet raw material alloy (A) of the present invention is an alloy mainly useful for the production method of the present invention described below, but its use is not necessarily limited to the production method of the present invention as long as it is for a rare-earth magnet. . The composition of the raw material alloy (A) is such that R 27.6 to 35.0 mass% of at least one selected from rare earth metal elements including yttrium, boron 0.94 to 1.30 mass%, and iron And the balance of M.
[0008]
The rare earth metal element of R is not particularly limited, but lanthanum, cerium, praseodymium, neodymium, yttrium, dysprosium and the like are preferable. If the content of R is less than 27.6% by mass, a large amount of α-Fe phase precipitates and the magnetic properties deteriorate, and if it exceeds 35.0% by mass, the ratio of the R-rich phase inside the sintered body is high. And the corrosion resistance is reduced.
If the boron content is less than 0.94% by mass, R₂Fe₁₇Phase precipitates and R₂Fe₁₄When the proportion of the B phase decreases, the residual magnetic flux density decreases. When the proportion exceeds 1.30% by mass, the proportion of the B-rich phase increases, and both the magnetic properties and the corrosion resistance decrease.
The iron content in the balance M is usually at least 50% by mass, preferably at least 60% by mass. The remainder M can include, in addition to iron, transition metals such as cobalt, aluminum, chromium, manganese, magnesium, copper, tin, tungsten, niobium, and gallium, carbon, silicon, and two or more of these. Further, the raw material alloy (A) may further contain unavoidable impurities in industrial production such as oxygen and nitrogen.
[0009]
The structure of the raw material alloy (A) of the present invention is R₂Fe₁₄B phase as main phase, R phase₂Fe₁₄90% by volume or more of the B phase is composed of crystal grains having a particle diameter in the short axis direction of 0.1 to 50 μm and a particle diameter in the long axis direction of 30 to 500 μm. Raw material alloy (A) is R₂Fe₁₄The B phase contains 77% by volume or more, preferably 77 to 85% by volume, and the R-rich phase contains 15 to 23% by volume, preferably 18 to 23% by volume. The R-rich phase is the R-rich phase.₂Fe₁₄It is preferable that the fine particles are finely dispersed so as to surround crystal grains having the B phase as a main phase. Part of the R-rich phase is R₂Fe₁₄It exists finely dispersed in the B phase crystal grains. When the content of the R-rich phase is less than 15% by volume, the effect of expanding the sintering temperature range is small when used in the production method of the present invention described below, and when it exceeds 23% by volume, the residual magnetic flux density decreases.
The particle diameters of the crystal grains in the short axis direction and the long axis direction of the raw material alloy (A) can be measured from a photograph of a cross-sectional structure of the raw material alloy (A) such as a flake taken by an optical microscope. The volume fraction of each crystal phase can be determined by EPMA image analysis.
In the present invention, a Comp image of a cross section in the center in the thickness direction of the alloy slab is photographed using JXA8800 manufactured by JEOL Ltd. Analysis was performed to determine the area ratio of each phase, which was defined as the volume ratio. Since the value obtained here is a value obtained by analyzing the alloy cross section in two dimensions, the value may be different from the actual volume ratio.
[0010]
R of raw alloy (A)₂Fe₁₄The c-axis lattice constant of the B phase is larger than the equilibrium state described later by 0.3 to 1%, preferably 0.3 to 0.8%. The c-axis lattice constant was obtained by pulverizing the alloy to 25 μm or less and using a powder X-ray diffraction method. The raw material alloy (A) exhibits such physical properties because of R₂Fe₁₄This is presumably because R atoms penetrate into the crystal lattice of the B phase and the crystal lattice is distorted.
When the increase in the c-axis lattice constant is less than 0.3% as compared with the c-axis lattice constant in the equilibrium state, the volume ratio of the R-rich phase is small, and₂Fe₁₄Since few R atoms penetrate into the B-phase lattice, the R-rich phase and R₂Fe₁₄The R component discharged from the B phase is small, and the liquid phase described later is not sufficiently supplied. Therefore, the sinterability is reduced and the holding power is reduced. Further, when used in the production method of the present invention described below, the effect of expanding the sintering temperature range is small, and the variation in the magnetic properties and shrinkage of the obtained rare earth magnet cannot be reduced. On the other hand, it is difficult to produce an alloy whose c-axis lattice constant exceeds 1% as compared with the c-axis lattice constant in an equilibrium state.
[0011]
The alloy in the equilibrium state refers to an alloy obtained by heat-treating an alloy at 600 to 800 ° C. for 40 hours or more, and then gradually cooling the alloy at a cooling rate of 1 ° C./min or less.
[0012]
The raw material alloy (A) has a higher volume ratio of the R-rich phase than the raw material alloy cast by a conventional strip casting method, and the R-rich phase exists in a state of being very finely dispersed. The R-rich phase is considered to be in a state in which more Fe and B are present as compared to the R-rich phase of the raw material alloy cast by the conventional strip casting method. The raw material alloy (A) having such a composition and crystal structure can be used in the production method of the present invention described below, so that the R-rich phase and the R-rich phase₂Fe₁₄The R atoms penetrating into the crystal lattice of the B phase create a liquid phase like the grain boundary phase alloy that supplies the R-rich phase in the conventional two-alloy method. R-rich phase in which R is finely dispersed;₂Fe₁₄Since the liquid phase is uniformly discharged, only the required amount of the liquid phase is more uniformly supplied to the portion required for sintering as compared with the grain boundary phase alloy of the conventional two-alloy method. . Further, since the raw material alloy (A) has a smaller amount of R as compared with the grain boundary phase alloy in the two-alloy method, the R-rich phase, which is a non-magnetic phase, is not generated more than necessary, and the residual magnetic flux density does not decrease.
[0013]
The raw material alloy (A) is desirably a flake having a plate thickness of 0.03 to 2.00 mm, preferably 0.10 to 1.0 mm.
The raw material alloy (A) has an R₂Fe₁₄When the hydrogen storage amount of the B phase is₂Fe₁₄It is 3 to 22% larger than the hydrogen storage amount of phase B, preferably 12 to 22% larger. Further, the hydrogen storage amount of the R-rich phase is 10 to 60% smaller, preferably 10 to 30% smaller than the hydrogen storage amount of the R-rich phase in an equilibrium state. Hereinafter, + is described when the hydrogen storage amount is larger than the equilibrium state, and-is shown when the hydrogen storage amount is smaller than the equilibrium state.
Such a hydrogen storage property is based on the R of the raw material alloy (A).₂Fe₁₄This is considered to be due to the characteristics of the compositions of the B phase and the R-rich phase. R of raw alloy (A)₂Fe₁₄The B phase contains more R components than usual and has good hydrogen storage properties, and the R-rich phase contains components other than R (Fe, B, etc.) more than usual, and these components are relatively hydrogen storage properties with R. Is considered to be low, and as a result, although the volume fraction of the R-rich phase is high, the hydrogen storage capacity is relatively low.
R₂Fe₁₄When the hydrogen storage amount of the B phase is₂Fe₁₄When the hydrogen storage capacity of the B phase is smaller than + 3.0% and the hydrogen storage capacity of the R-rich phase is larger than the hydrogen storage capacity of the equilibrium R-rich phase by -10%, R₂Fe₁₄It is not preferable because the B phase becomes coarse, and the sinterability may be reduced and the coercive force may be reduced. R₂Fe₁₄When the hydrogen storage amount of the B phase is₂Fe₁₄If the hydrogen storage capacity of the B-phase is greater than + 22% and the hydrogen storage capacity of the R-rich phase is smaller than -60% as compared with the hydrogen storage capacity of the equilibrium R-rich phase, Is difficult to manufacture.
Here, the equilibrium state is as described above, and the hydrogen storage amount can be measured by a PCT device as follows.
[0014]
First, the alloy to be measured is put into a cell having a constant capacity, and the pressure is evacuated to 1 to 3 Pa. The temperature is kept at 40 ° C. in a reduced pressure atmosphere, and a hydrogen atmosphere of 4 atm is set. Next, the hydrogen pressure at the time when the hydrogen absorption of the alloy is stopped is measured. At this time, if the hydrogen pressure at the time of stopping the hydrogen storage is 1 atm or less, the hydrogen atmosphere is again placed at 4 atm. This is repeated until the pressure at the time of stopping hydrogen storage becomes 1 atm or more. The hydrogen storage amount of the alloy can be determined from the sum of these hydrogen pressure fluctuations, and the hydrogen storage amount per unit weight can be determined from the weight of the alloy charged into the cell.
Next, the above measurement is also performed on alloy samples obtained under the same casting conditions and having different amounts of R addition, and a graph of the R addition amount on the horizontal axis and the hydrogen storage amount on the vertical axis is created. From the approximate curve of the graph, R₂Fe₁₄The hydrogen storage amount at the R amount (26.7% by mass) of the stoichiometric composition of the B phase was determined, and this was calculated as R₂Fe₁₄The hydrogen storage amount of B phase. Hydrogen storage amount and R of the whole alloy in each composition₂Fe₁₄The hydrogen storage amount of the R-rich phase can be determined from the difference from the hydrogen storage amount of the B phase.
R of equilibrium alloy₂Fe₁₄The hydrogen storage amount of the B phase and the hydrogen storage amount of the R-rich phase can be determined by the same method.
[0015]
In the raw material alloy (A), the volume ratio (Rx) of the R-rich phase is represented by Rx = Rc × X (where Rc is the volume ratio of the R-rich phase in an equilibrium state). Is preferably 2.2 to 5.0, particularly preferably 3.0 to 5.0. The volume ratio (Rx) of the R-rich phase and the volume ratio (Rc) of the R-rich phase in an equilibrium state can be determined by EPMA image analysis.
When X exceeds 5.0, the volume ratio of the R-rich phase increases, and the residual magnetic flux density of the obtained magnet may decrease. When X is less than 2.2, the liquid phase is sufficiently supplied as described above. This is not preferable because the sinterability may decrease and the coercive force may decrease.
[0016]
The raw alloy powder for a rare earth magnet of the present invention is a powder obtained by subjecting the above-described raw material alloy (A) of the present invention to hydrogen absorption and collapse, and is an alloy powder useful for the production method of the present invention described below. If it is for a rare earth magnet, its use is not necessarily limited to the production method of the present invention.
The raw material alloy powder for a rare earth magnet of the present invention is a powder that contains 50% by mass or more of a powder having a particle size of 355 to 850 µm and a weight average particle size of 300 to 800 µm, preferably 300 to 500 µm.
The hydrogen storage decay is performed, for example, by allowing the raw material alloy (A) to store hydrogen for 30 minutes or more in an atmosphere at a hydrogen pressure of 0.1 to 0.4 MPa and a normal temperature to 100 ° C. The evacuation can be performed by evacuation until an atmosphere is obtained, and the like.
The raw alloy powder for a rare earth magnet of the present invention obtained by hydrogen absorption and disintegration has a small amount of extremely fine powder, suppresses oxidation, has a good yield of R, and has a very small amount of extremely coarse pulverized powder. The pulverizing step can be performed efficiently.
[0017]
The raw material alloy (A) of the present invention can be prepared, for example, by the following method.
First, a raw material metal or a master alloy of R, boron, and M adjusted to the composition of the above-mentioned raw material alloy (A) is melted by a high-frequency melting method in an inert gas atmosphere, and the molten material is rolled into a single roll or twin roll. It is continuously solidified by a strip casting method using a roll or a disk. At that time, the average cooling rate from the melting point of the alloy to the roll peeling can be performed at 50 to 3000 ° C./sec (primary cooling step). By the primary cooling step, R₂Fe₁₄The particle diameter of the B phase in the short axis direction and the long axis direction, R₂Fe₁₄Volume ratio of B phase and R-rich phase, R₂Fe₁₄The approximate c-axis lattice constant of the B phase is determined.
If the average cooling rate is too high, chill crystals are generated, and the particle diameters R and R in the specified short axis direction and long axis direction are determined.₂Fe₁₄Crystal grains having the B phase as the main phase cannot have a prescribed volume ratio, and the residual magnetic flux density decreases. If it is too late, R₂Fe₁₄The c-axis lattice constant of the B phase does not fall within the specified range, and R₂Fe₁₄The B phase coarsens, the dispersibility of the R-rich phase decreases, and the coercive force decreases.
[0018]
Further, by controlling the cooling process of the slab after the roll is peeled, R₂Fe₁₄An alloy in which the c-axis lattice constant and hydrogen storage amount of the B phase are preferably controlled can be uniformly produced. For example, after the primary cooling step, the slab is cooled to a temperature of ternary eutectic point + 30 ° C. at an average cooling rate of 30 ° C./sec or more (secondary cooling step), and a ternary eutectic temperature of ± 30 ° C. A method of cooling the range within 30 seconds (a tertiary cooling step) is exemplified. The tertiary cooling step is preferably performed within 5 seconds. This step can be preferably performed to control the precipitated phase in the R-rich phase to control the hydrogen storage characteristics. In the third cooling step, if the holding time in the range of ternary eutectic temperature ± 30 ° C. is too long, R₂Fe₁₄There is a possibility that the B phase becomes coarse, segregation of the R-rich phase occurs, and the coercive force decreases. Thereafter, the raw material alloy (A) of the present invention can be obtained by cooling to preferably 100 ° C. or lower at 5 to 30 ° C./min (fourth cooling step).
The temperature control in the production of the raw material alloy (A) can be appropriately performed by selecting a member of a casting apparatus to be brought into contact with the cast slab after the roll is detached, a temperature control apparatus having a heating mechanism or a cooling mechanism, or the like.
[0019]
The method for producing a rare earth magnet according to the present invention is characterized in that the raw material alloy (A) and R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements including yttrium, and boron 0.94 to 1 30% by mass and the balance M containing iron, and R as a main phase₂Fe₁₄A preparation step of preparing a raw material alloy (B) having a B phase content of 85% by volume or more and an R-rich phase content of 15% by volume or less, and pulverizing the material alloys (A) and (B) The method includes a pulverizing step, a molding step of molding the obtained pulverized powder, a sintering step of sintering the obtained molded article, and a heat treatment step of aging the obtained sintered article.
[0020]
As the material alloy (A) used in the preparation step, the above-described material alloy (A) can be used, and as the material alloy (A), the above-described material alloy powder for a rare earth magnet of the present invention can also be used.
The composition range of the raw material alloy (B) used in the preparation step is the same as that of the raw material alloy (A), and the same elements as the raw material alloy (A) described above can be used for the elements that may be contained as necessary. Therefore, the compositions of the raw material alloys (A) and (B) may be the same or different in the specified composition range.
[0021]
R of raw material alloy (B)₂Fe₁₄It is sufficient that the volume ratio of the B phase is 85% or more, and the volume ratio can be obtained by the same method as described above.
The crystal structure of the material alloy (B) may be the same or different except that the content ratio of the R-rich phase is different from that of the material alloy (A). The content ratio of the R-rich phase in the raw material alloy (B) is 15% by volume or less, preferably 3 to 10% by volume. When the content ratio of the R-rich phase of the raw material alloy (B) exceeds 15% by volume, the residual magnetic flux density decreases.
[0022]
As described above, the raw material alloy (A) prepared in the preparation step supplies the liquid phase under the optimum condition,₂Fe₁₄B phase is also supplied. On the other hand, the raw material alloy (B) is mainly composed of R, like the main phase alloy in the conventional two-alloy method.₂Fe₁₄It plays the role of supplying the B phase. Since the raw alloys (A) and (B) do not use alloys having a high R content unlike the conventional two-alloy method, oxidation of R is suppressed, and the yield of R is good. The effect of the raw material alloy (A) can widen the optimum sintering temperature range as compared with the case where a single alloy is used, so that variations in the magnetic properties and shrinkage of the obtained rare earth magnet can be suppressed.
[0023]
In the preparation step, the mixing ratio of the raw material alloys (A) and (B) can be appropriately selected according to the magnetic properties of the rare earth magnet, the characteristics of the raw material alloy (A), and the characteristics of the raw material alloy (B). The ratio of the raw material alloy (A) :( B) = 1: 1 to 30 is preferable. As the material alloy (B), those having an appropriate composition or crystal structure within a specified range can be used. For example, when obtaining a permanent magnet having a high residual magnetic flux density, it is preferable to use an alloy having an R content of 28.1 to 30.0% by mass as the raw material alloy (B). With such an R content, the volume fraction of the main phase is increased, and the residual magnetic flux density is improved. When an alloy having an R content of 30.0 to 33.0% by mass is used, R₂Fe₁₄If the average particle diameter in the minor axis direction of the B phase is 5 μm or more, the ratio of two or more main phases having different crystal orientations in the pulverized powder is reduced when pulverized in the manufacturing process. Is obtained.
When a permanent magnet having a high coercive force is obtained, it is preferable to use an alloy having an R content of 31 to 35% by mass as the raw material alloy (B). With such an R content, as the R content increases, the volume ratio of the R-rich phase increases, the sinterability improves, and the coercive force improves. Even if the R content of the raw material alloy (B) is less than 31% by mass, R₂Fe₁₄When the average particle size of the B phase in the uniaxial direction is 4.0 μm or less and the content ratio of the R-rich phase is about 15% by volume, the sinterability is improved and a high coercive force is obtained. When such a raw material alloy (B) is used, the mixing ratio of the raw material alloy (A) is preferably 50% by mass or less based on all the raw material alloys.
[0024]
In the preparation step, the raw material alloy (A) can be obtained by the above-described method or the like, while the raw material alloy (B) can also be obtained by appropriately changing the manufacturing conditions. For example, it can be obtained by changing the composition or controlling the cooling rate including the temperature raising and holding steps in the cooling step.
The raw material alloys (A) and (B) may be in the form of a flake obtained by a strip casting method, or in the form of a coarsely pulverized powder obtained by coarsely pulverizing a flake. The raw material alloy powder for a rare earth magnet of the present invention is exemplified. Further, as long as the raw material alloy (B) falls within the above-specified range, it may be in the form of coarsely pulverized powder obtained by coarsely pulverizing an ingot obtained by die casting.
[0025]
Next, in the production method of the present invention, a pulverizing step is performed. In the pulverization step, usually, after hydrogenation pulverization, the raw material alloys (A) and (B) are pulverized to an average particle size of about 3 to 6 μm using a pulverizer such as a jet mill. The pulverization is preferably performed in a state where the raw material alloys (A) and (B) are mixed from the viewpoint of workability or suppressing oxidation. However, it is also possible to mix the crushed powders after crushing each.
[0026]
In the production method of the present invention, the obtained pulverized powder is then molded into a desired shape and size. The molding can be performed by a known method employed in the manufacture of rare earth magnets, for example, by a method of molding by applying pressure in a magnetic field. Usually, 0.5 to 3.0 t / cm in a magnetic field of 15 to 30 kOe.²Mold under pressure.
[0027]
In the manufacturing method of the present invention, a sintering step for sintering the molded product is performed. The sintering temperature in the sintering step can be appropriately selected from the range of usually 1000 to 1100 ° C, and the sintering time can also be appropriately selected from the range of usually 1 to 5 hours. The sintering temperature range in the above uses the above-mentioned raw material alloy (A), so that the magnetic characteristics and shrinkage of the rare earth magnet can be obtained even when set in a wider range as compared with the case of using a conventional single alloy. Variation can be reduced. Therefore, the sintering temperature control is eased as compared with the case of using the conventional single alloy.
[0028]
In the manufacturing method of the present invention, a heat treatment step of aging the sintered product is performed. The aging treatment can also be performed by appropriately selecting conditions from known methods to obtain a desired rare earth magnet. The aging treatment can be performed by, for example, a method of lowering the temperature in a temperature range of 450 to 950 ° C. twice or more and maintaining the temperature for a desired time.
[0029]
【Example】
Hereinafter, the present invention will be described in more detail with reference to Examples and Comparative Examples, but the present invention is not limited thereto.
Example 1
(Preparation of raw material alloy (A))
Nd 32.0% by mass, Dy 1.0% by mass, B 1.00% by mass, Al 0.20% by mass, Co 1.0% by mass, Neodymium metal, Dysprosium metal, Ferroboron, Aluminum, cobalt, and iron were blended and melted in a high-frequency melting furnace using an alumina crucible in an argon gas atmosphere. Next, the obtained molten alloy was cast by a strip casting method using a water-cooled copper single-roll casting apparatus to obtain a slab having a thickness of about 0.2 mm. The ternary eutectic point of this alloy is about 640 ° C. The temperature of the molten metal immediately before coming into contact with the roll of the alloy was about 1350 to 1400 ° C, and the temperature of the slab immediately after peeling off the roll was about 600 ° C when measured with an infrared thermal image measurement device. The cooling time on the roll was about 0.6 seconds.
Next, the cast slab after the roll peeling was cooled by a rotating drum type water cooling device, and recovered after 40 minutes to obtain Sample 1 as a raw material alloy (A). The time required to enter the water cooling device after the roll was peeled was about 0.8 seconds. The slab temperature immediately after the slab was moved to the water cooling device was about 450 ° C., and the slab temperature immediately after the removal was about 60 ° C.
[0030]
The cross-sectional structure of the obtained cast slab (sample 1) was photographed with an optical microscope.₂Fe₁₄The average particle diameter of the B phase in the short axis direction and the average particle diameter in the long axis direction were measured, and the respective average particle diameters were determined.₂Fe₁₄The volume ratio of the B phase and the volume ratio of the R-rich phase were determined. The volume fraction of crystal grains (X) having an average particle diameter in the short axis direction of 0.1 to 50 μm and an average particle diameter in the long axis direction of 30 to 500 μm was determined. The particle diameter in the short axis direction is 3.3 μm, the particle diameter in the long axis direction is 74 μm, R₂Fe₁₄The volume ratio of the B phase was 82% by volume, the volume ratio of the crystal grains (X) was 95% by volume, and the volume ratio of the R-rich phase was 18% by volume. After the slab was ground to about 25 μm,₂Fe₁₄The c-axis lattice constant of the B phase was measured and found to be 12.34 °.
The obtained slab is hydrogenated in a hydrogen atmosphere of 30 ° C. and 0.1 MPa for 1 hour, and then dehydrogenated at 400 ° C. to perform hydrogen pulverization, and the obtained pulverized powder is subjected to a low tap standard sieve shaker. As a result, the powder having a particle diameter of 355 to 850 μm was about 74% by mass, and the average particle diameter of the powder was about 450 μm. Further, the amount of hydrogen occlusion determined by a PCT apparatus was 0.393% by mass. From the amount of occlusion, R₂Fe₁₄When the hydrogen storage amounts of the B phase and the R-rich phase were determined, about 0.278% by mass was occluded in the main phase and 0.115% by mass was occluded in the R-rich phase.
The obtained slab was heat-treated at 800 ° C. for 40 hours to be in an equilibrium state, cooled at a cooling rate of 1 ° C./min or less, and analyzed by the above-mentioned method.₂Fe₁₄The c-axis lattice constant of the B phase is 12.25 °, the hydrogen absorption is 0.408% by mass, the hydrogen absorption of the main phase is 0.237% by mass, the hydrogen absorption of the R-rich phase is 0.171% by mass, The volume fraction of the R-rich phase was about 4%.
Tables 1 and 2 show the above measurement results.
[0031]
(Adjustment of raw material alloy (B))
Nd 32.0% by mass, Dy 1.0% by mass, B 1.00% by mass, Al 0.20% by mass, Co 1.0% by mass, and the balance iron so that neodymium metal, dysprosium metal, ferroboron, Aluminum, cobalt, and iron were blended and melted in an argon gas atmosphere using an alumina crucible in a high-frequency melting furnace. Next, it was cast by a strip casting method using a water-cooled copper single-roll casting apparatus to obtain a slab having a thickness of about 0.4 mm. The temperature of the molten metal immediately before coming into contact with the roll was 1300 to 1350 ° C, and the temperature of the slab immediately after peeling from the roll was about 800 ° C when measured with an infrared thermal image measurement device. The cooling time on the roll was about 1.2 seconds.
Next, the slab after the roll peeling was collected in a steel container, the container was sealed, taken out into the atmosphere and allowed to cool, and collected after 1500 minutes to obtain a sample 1a as a material alloy (B). . The temperature of the slab in the container was about 665 ° C. immediately after the collection, about 615 ° C. after 80 minutes from the collection, and about 90 ° C. when the slab was taken out after 1500 minutes. Also, the time required to enter the steel container after the roll was peeled was about 0.8 seconds. Sample 1a was analyzed in the same manner as Sample 1 described above. The results are shown in Tables 1 and 2. FIG. 1 is a graph showing the thermal history at the time of manufacturing the samples 1 and 1a.
[0032]
(Method of manufacturing permanent magnets)
Sample 1 and sample 1a prepared above were introduced into a drum mixer at a mass ratio of 5: 5 and mixed. After hydrogenation at 30 ° C. and a hydrogen atmosphere of 0.1 MPa for 1 hour, hydrogen pulverization was performed by dehydrogenation at 400 ° C. Next, pulverization was performed by a jet mill so that the average particle diameter became 5.0 μm.
Then, in a magnetic field of 15 kOe, 2.5 ton / cm², And the obtained molded body was sintered in a vacuum for 4 hours. At that time, the sintering temperature was changed to 1055 ° C, 1060 ° C and 1065 ° C. After sintering, the first stage heat treatment was performed at 900 ° C. for 1 hour, and the second stage heat treatment was performed at 500 ° C. for 2 hours to perform aging treatment. The magnetic properties of the obtained permanent magnet were measured by a conventional method. Table 3 shows the results.
The orientation shrinkage of the obtained permanent magnet was measured according to the following definition. Generally, in the case of an R—Fe—B based sintered magnet, press molding is performed while orienting particles in a magnetic field in order to magnetically anisotropy. Accordingly, the amount of shrinkage during sintering differs between the particle orientation direction (c-axis direction) and the a-axis direction perpendicular thereto. ΔL for shrinkage in orientation direction, sintering and aging
The length of the molded body before processing is L₀In this case, the orientation shrinkage can be determined by the following equation. Table 3 shows the results.
Shrinkage = ΔL / L₀
[0033]
Example 2 ~ 4
A permanent magnet was prepared in the same manner as in Example 1 except that the mixing ratio of Sample 1 and Sample 1a prepared in Example 1 was changed as shown in Table 2, and each magnetic property and the like were measured. Table 3 shows the results.
[0034]
Example 5
Example 1 except that the alloy composition was 34.0% by mass of Nd, 1.0% by mass of Dy, 1.00% by mass of B, 0.20% by mass of Al, 1.0% by mass of Co and the balance iron. In the same manner as in Sample 1, a sample 2 as a raw material alloy (A) was prepared. The ternary eutectic point of this alloy is about 640 ° C. The heat history of the slab during casting was almost the same as that of Sample 1.
Also, except that the composition of the alloy was 31.5% by mass of Nd, 1.0% by mass of Dy, 1.0% by mass of B, 0.20% by mass of Al, 1.0% by mass of Co and the balance iron. A sample 2a, which is a raw material alloy (B), was prepared in the same manner as the sample 1a in Example 1. The ternary eutectic point of this alloy is about 640 ° C. The heat history of the slab during casting was almost the same as that of sample 1a. The same analysis as Sample 1 in Example 1 was performed on the obtained Sample 2 and Sample 2a. The results are shown in Tables 1 and 2.
Further, a permanent magnet was prepared in the same manner as in Example 1 except that Sample 2 and Sample 2a were mixed at a mass ratio of 2: 8, and the magnetic properties and the like were measured. Table 3 shows the results.
[0035]
Comparative example 1
A permanent magnet was prepared in the same manner as in Example 1 except that only the sample 1a prepared in Example 1 was used, and the magnetic properties and the like were measured. Table 3 shows the results.
[0036]
Comparative example 2
Except that the alloy composition was 30.5% by mass of Nd, 1.11% by mass of B, 0.20% by mass of Al, and the balance iron, the raw material alloy (B) was similar to the sample 1a in Example 1. Was prepared as sample 3a. The ternary eutectic point of this alloy is about 640 ° C. The heat history of the slab during casting was almost the same as that of sample 1a.
Also, except that the alloy composition was Nd 45.5% by mass, Dy 10.0% by mass, Al 0.20% by mass, Co 10.0% by mass, and the balance iron, and the heat history was as follows. A sample 4a different from the material alloys (A) and (B) was obtained in the same manner as the sample 1a in Example 1. At this time, the temperature of the molten metal immediately before coming into contact with the roll was about 1450 to 1500 ° C, and the temperature of the cast piece immediately after being peeled off the roll was about 550 ° C. The cooling time on the roll was about 1.8 seconds. Further, the slab after the roll peeling was cooled by a rotating drum type water cooling device, and was recovered after a lapse of 40 minutes. The slab temperature immediately after the slab was moved to the water cooling device was about 400 ° C., and the slab temperature immediately after the removal was about 55 ° C. The slab thickness of each of Sample 3a and Sample 4a was about 0.4 mm.
For sample 3a, the same analysis as for sample 1 in Example 1 was performed. The results are shown in Tables 1 and 2.
Further, a permanent magnet was prepared in the same manner as in Example 1 except that the obtained sample 3a and sample 4a were mixed at a mass ratio of 9: 1, and magnetic properties and the like were measured. Table 3 shows the results.
[0037]
[Table 1]

[0038]
[Table 2]

[0039]
[Table 3]

[0040]
【The invention's effect】
In the method for producing a rare earth magnet of the present invention, particularly, the raw material alloy (A) of the present invention or the rare earth magnet raw material alloy of the present invention is used in combination with the raw material alloy (B). Therefore, the conventional method of manufacturing a rare earth magnet using a single alloy is alleviated, and the problem of oxidation of the R component and the poor yield in the conventional two alloy method are improved. Moreover, the obtained rare-earth magnet does not have a decrease in magnetic properties due to a variation in sintering temperature and a variation in shrinkage, and is superior in magnetic properties to a permanent magnet prepared by a conventional two-alloy method.
Further, since the raw material alloy (A) for rare earth magnets and the powder thereof of the present invention have a specific composition and structure, they are useful for a production method for stably and easily obtaining rare earth magnets having high magnetic properties. It is useful for the production method of the present invention.
[Brief description of the drawings]
FIG.
3 is a graph showing the thermal history of Sample 1 and Sample 1a prepared in Example 1 during manufacture.

Claims

R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements containing yttrium, 0.94 to 1.30% by mass of boron, and a balance of iron-containing balance M , R ₂ Fe ₁₄ B phase as the main phase, the content ratio of the R ₂ Fe ₁₄ B phase is 77% by volume or more, the content ratio of the R-rich phase is 15 to 23% by volume, and the R ₂ Fe ₁₄ B phase 90% by volume or more is composed of crystal grains having a particle diameter in the short axis direction of 0.1 to 50 μm and a particle diameter in the long axis direction of 30 to 500 μm, and the c-axis lattice constant of the R ₂ Fe ₁₄ B phase is in an equilibrium state. 0.3 to 1% larger raw material alloy (A) and
R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements containing yttrium, 0.94 to 1.30% by mass of boron, and a balance of iron-containing balance M A preparation step of preparing a raw material alloy (B) in which the content of the R ₂ Fe ₁₄ B phase as the main phase is 85% by volume or more and the content of the R-rich phase is less than 15% by volume,
A crushing step of crushing the raw material alloys (A) and (B);
A molding step of molding the obtained ground powder,
A sintering step of sintering the obtained molded product,
And a heat treatment step of heat-treating the obtained sintered product.

R 27.6 to 35.0% by mass of at least one selected from rare earth metal elements containing yttrium, 0.94 to 1.30% by mass of boron, and a balance of iron-containing balance M , R ₂ Fe ₁₄ B phase as the main phase, the content ratio of the R ₂ Fe ₁₄ B phase is 77% by volume or more, the content ratio of the R-rich phase is 15 to 23% by volume, and the R ₂ Fe ₁₄ B phase 90% by volume or more is composed of crystal grains having a particle diameter in the short axis direction of 0.1 to 50 μm and a particle diameter in the long axis direction of 30 to 500 μm, and the c-axis lattice constant of the R ₂ Fe ₁₄ B phase is in an equilibrium state. A raw material alloy for rare earth magnets (A), which is 0.3 to 1% larger than the alloy.

The raw material alloy (A) according to claim 2, wherein the balance M includes at least one selected from the group consisting of transition metal elements other than iron, silicon, and carbon.

The raw material alloy (A) according to claim 2 or 3, wherein the form is a flake having a thickness of 0.03 to 2.0 mm.

Hydrogen storage capacity of the _R 2 _{Fe 14} B phase per alloy unit mass, 3.0 to 22% than the hydrogen storage capacity of the _R 2 _{Fe 14} B phase in equilibrium increases, hydrogen R-rich phase per alloy unit mass The raw material alloy (A) according to any one of claims 2 to 4, wherein the storage amount is 10 to 60% smaller than the hydrogen storage amount of the R-rich phase in an equilibrium state.

When the volume ratio (Rx) of the R-rich phase is expressed as Rx = Rc × X (where Rc is the volume ratio of the R-rich phase in an equilibrium state), X is 2.2 to 5.x. The raw material alloy (A) according to any one of claims 1 to 5, which is 0.

A powder obtained by subjecting the raw material alloy (A) according to any one of claims 2 to 6 to hydrogen absorption and disintegration, comprising a powder having a particle size of 355 to 850 µm in an amount of 50% by mass or more, and having a weight average particle size of 300. A raw material alloy powder for rare earth magnets, which has a thickness of from 800 to 800 µm.