JP2005320628A - Alloy slug for r-t-b-based sintered magnet, manufacturing method therefor and magnet - Google Patents
Alloy slug for r-t-b-based sintered magnet, manufacturing method therefor and magnet Download PDFInfo
- Publication number
- JP2005320628A JP2005320628A JP2005110368A JP2005110368A JP2005320628A JP 2005320628 A JP2005320628 A JP 2005320628A JP 2005110368 A JP2005110368 A JP 2005110368A JP 2005110368 A JP2005110368 A JP 2005110368A JP 2005320628 A JP2005320628 A JP 2005320628A
- Authority
- JP
- Japan
- Prior art keywords
- alloy
- length
- axis direction
- rich phase
- minor axis
- Prior art date
- Legal status (The legal status is an assumption and is not a legal conclusion. Google has not performed a legal analysis and makes no representation as to the accuracy of the status listed.)
- Granted
Links
- 0 C(*1)C11C=C**1 Chemical compound C(*1)C11C=C**1 0.000 description 1
Images
Abstract
Description
本発明は、希土類合金、特にＲ−Ｔ−Ｂ系焼結磁石用合金塊、その製造法及び該合金塊を用いた磁石に関する。 The present invention relates to a rare earth alloy, particularly an alloy block for an RTB-based sintered magnet, a manufacturing method thereof, and a magnet using the alloy block.
近年、磁石用合金としてＮｄ−Ｆｅ−Ｂ系合金がその高特性から急激に生産量を伸ばしており、ＨＤ（ハードディスク）用、ＭＲＩ（磁気共鳴映像法）用あるいは、各種モーター用等に使用されている。通常は、Ｎｄの一部をＰｒ、Ｄｙ等他の希土類元素で置換したもの（Ｒと表記する。）、またＦｅの一部をＣｏ、Ｎｉ等他の遷移元素で置換したもの（Ｔと表記する。）が一般的であり、Ｎｄ−Ｆｅ−Ｂ系合金を含め、Ｒ−Ｔ−Ｂ系合金と総称されている。 In recent years, Nd-Fe-B alloys as magnet alloys have rapidly increased in production due to their high characteristics, and are used for HD (hard disk), MRI (magnetic resonance imaging) or various motors. ing. Usually, a part of Nd is substituted with other rare earth elements such as Pr and Dy (denoted as R), and a part of Fe is substituted with other transition elements such as Co and Ni (denoted as T) Are generally called R-T-B alloys including Nd-Fe-B alloys.
Ｒ−Ｔ−Ｂ系合金は、磁化作用に寄与する強磁性相Ｒ2Ｔ14Ｂを主相とする結晶と、非磁性で希土類元素の濃縮した低融点のＲ−リッチ相を結晶粒界に持つ合金で、活性な金属であることから一般に真空又は不活性ガス中にて溶解され、金型に鋳造されてきた。
この合金塊は、粉砕され３μｍ（ＦＳＳＳ：フィッシャーサブシーブサイザーでの測定）程度の粉体とした後、磁場中でプレス成形され、焼結炉で約１０００〜１１００℃の高温にて焼結され、その後必要に応じ熱処理、機械加工され、耐食のためのメッキをされ磁石化されるのが普通である。
In the R-T-B system alloy, a crystal having a ferromagnetic phase R 2 T 14 B that contributes to a magnetization action as a main phase and a low-melting R-rich phase that is nonmagnetic and enriched with rare earth elements is used as a grain boundary. Since it is an active metal, it is generally melted in a vacuum or an inert gas and cast into a mold.
This alloy lump is pulverized into a powder of about 3 μm (measured with FSSS: Fischer Sub-Seeb Sizer), press-molded in a magnetic field, and sintered at a high temperature of about 1000 to 1100 ° C. in a sintering furnace. Then, it is usually heat treated and machined as necessary, plated for corrosion resistance and magnetized.
このＲ−リッチ相は、以下の点で重要な役割を担っている。
１）融点が低く、焼結時に液相となり、磁石の高密度化、従って磁化の向上に寄与する。
２）粒界の凹凸を無くし、逆磁区のニュークリエーションサイトを減少させ保磁力を高める。
３）主相を磁気的に絶縁することから保磁力を高める。
従ってＲ−リッチ相の分散状態が悪いと磁石としての特性に影響するため、均一であることが重要となる。
最終的な磁石としてのＲ―リッチ相の分布は、原料用合金塊の組織に大きく影響される。すなわち、金型にて鋳造された場合、冷却速度が遅いため往々にして結晶粒が大きくなる。この結果、粉砕した時の粒が結晶粒径よりはるかに細かくなり、金型鋳造ではＲ−リッチ相はほとんどが結晶粒界に凝集し粒内に無いため、Ｒ−リッチ相を含まない主相のみの粒とＲ−リッチ相のみの粒とが別々に存在し均一な混合がしにくくなる。
This R-rich phase plays an important role in the following points.
1) The melting point is low and it becomes a liquid phase at the time of sintering, which contributes to increasing the density of the magnet and thus improving the magnetization.
2) Eliminate grain boundary irregularities, reduce reverse domain nucleation sites and increase coercivity.
3) Since the main phase is magnetically insulated, the coercive force is increased.
Therefore, if the dispersion state of the R-rich phase is poor, it affects the characteristics of the magnet, and therefore it is important that it is uniform.
The distribution of the R-rich phase as the final magnet is greatly influenced by the structure of the raw material alloy ingot. That is, when cast in a mold, the cooling rate is slow and the crystal grains are often large. As a result, the grains when pulverized become much finer than the crystal grain size, and in the die casting, the R-rich phase is mostly aggregated at the grain boundaries and is not in the grains, so that the main phase does not contain the R-rich phase. Only particles and only R-rich phase particles are present separately, making uniform mixing difficult.
金型鋳造でのもう一つの問題は、冷却速度が遅いため初晶としてγ―Ｆｅが生成しやすくなることである。γ―Ｆｅは約９１０℃以下では、α―Ｆｅに変態する。この変態したα―Ｆｅは、磁石製造時の粉砕効率の悪化をもたらし、焼結後も残存すれば磁気特性の低下をもたらす。そこで金型にて鋳造したインゴットの場合は、高温で長時間にわたる均質化処理によるα―Ｆｅの消去が必要となってくる。 Another problem in mold casting is that γ-Fe is likely to be formed as primary crystals due to the slow cooling rate. γ-Fe transforms into α-Fe at about 910 ° C or less. This transformed α-Fe deteriorates the pulverization efficiency at the time of magnet production, and if it remains after sintering, it causes a decrease in magnetic properties. Therefore, in the case of an ingot cast by a mold, it is necessary to erase α-Fe by a homogenization treatment for a long time at a high temperature.
これらを解決するため、金型鋳造方法より速い冷却速度で鋳造する方法として、ストリップキャスティング法（ＳＣ法と略す。）が紹介され実際の工程にて使用されている。
これは内部が水冷された銅ロール上に溶湯を流し、０．３ｍｍ程度の薄帯を鋳造することにより、急冷凝固させるものであり、結晶組織を微細化させ、Ｒ−リッチ相が微細に分散した組織を有する合金片を生成させるものであり、合金片内の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。（特許文献１及び２参照）しかし、この方法においてもＲ成分の濃度が低下するに従ってα−Ｆｅの発生は避けがたく、例えばＮｄ−Ｆｅ−Ｂの３元合金では、Ｎｄが２８質量％以下では、α−Ｆｅの発生が見られるようになる。
このα−Ｆｅは、磁石製造工程において粉砕性を著しく阻害する。
In order to solve these problems, a strip casting method (abbreviated as SC method) has been introduced and used in actual processes as a method of casting at a faster cooling rate than the mold casting method.
This is a method in which a molten metal is poured onto a copper roll that has been internally cooled with water, and a thin strip of about 0.3 mm is cast to solidify rapidly, thereby refining the crystal structure and finely dispersing the R-rich phase. The alloy piece having the texture is produced, and the R-rich phase in the alloy piece is finely dispersed, so the dispersibility of the R-rich phase after pulverization and sintering is improved, and the magnetic properties are improved. Has succeeded. (See Patent Documents 1 and 2) However, even in this method, the generation of α-Fe is unavoidable as the concentration of the R component decreases. For example, in a ternary alloy of Nd—Fe—B, Nd is 28% by mass or less. Then, generation | occurrence | production of (alpha) -Fe comes to be seen.
This α-Fe significantly impairs grindability in the magnet manufacturing process.
本発明者らは、従来の遠心鋳造法を改良し回転する鋳型の内側に配置した、往復運動し複数のノズルを備えた箱型のタンディッシュを介して、溶湯を回転鋳型の内側に堆積凝固させる方法（Ｃｅｎｔｒｉｆｕｇａｌ Ｃａｓｔｉｎｇ：以下ＣＣ法と略す。）と装置を発明した（特許文献３及び４参照）。 The present inventors improved the conventional centrifugal casting method, and deposited and solidified the molten metal on the inner side of the rotating mold through a box-type tundish that was reciprocated and provided with a plurality of nozzles, which was disposed inside the rotating mold. (Centrifugal Casting: hereinafter abbreviated as CC method) and an apparatus were invented (see Patent Documents 3 and 4).
ＣＣ法では既に堆積凝固した合金塊の上に次の溶湯が順次注がれ、追加鋳造されたその溶湯は鋳型が１回転する間に凝固するため、凝固速度を速めることができる。しかし、このＣＣ法でもＲ成分の濃度の低い合金を製造しようとすると、高温域の冷却速度が遅いためα−Ｆｅの生成は避けられない。 In the CC method, the next molten metal is sequentially poured onto the alloy mass that has already been deposited and solidified, and the additionally cast molten metal is solidified during one rotation of the mold, so that the solidification rate can be increased. However, even in this CC method, when an alloy having a low concentration of the R component is to be produced, the formation of α-Fe is inevitable because the cooling rate in the high temperature range is slow.
α−Ｆｅの生成を避けるため、本発明者らはＣＣ法で凝固冷却速度を速めるため溶湯の堆積速度をより小さくするため回転するタンディッシュから溶湯を飛散させ回転鋳型へ堆積させる遠心鋳造方法を考案した（Ｎｅｗ Ｃｅｎｔｒｉｆｕｇａｌ Ｃａｓｔｉｎｇ：以下ＮＣＣ法と略す。特許文献５参照）。これによりα−Ｆｅの発生が抑制されることがわかった。また、磁石として磁化特性を上げるための手段としてＲ成分の濃度が低い側でα−Ｆｅを実質的に含まない鋳造塊が得られるようになった。さらにＲ−リッチ相を微細均一な分布にするため、内面が凹又は／及び凸状の非平滑面をもつ回転する円筒状鋳型の内面に堆積凝固させる方法が考案された（特許文献６参照）。
さらに円筒状鋳型の内面に該鋳型の材質より熱伝導率が小さい膜が設けられている鋳型を用いて堆積凝固させる方法が考案された（特許文献７参照）。
Furthermore, a method of depositing and solidifying using a mold in which a film having a lower thermal conductivity than the material of the mold is provided on the inner surface of the cylindrical mold has been devised (see Patent Document 7).
特許文献６の方法により、Ｒ−リッチ相の分散性は向上したが、溶湯液滴を堆積させている際に、堆積済み合金塊の温度が上昇しＲ−リッチ相が凝集してプール状になり、このため焼結磁石製造工程の微粉砕工程においてＲ−リッチ相が先に粉砕されてしまい、得られる粉体組成の時間変動が安定しないという問題が起きた。また、得られた粉体のＲ−リッチ相の分散性もＳＣ法により鋳造された鋳片（以下ＳＣ合金と略す。）よりも悪く、このため保磁力も低めであるという問題も起きた。
また特許文献７の方法では、冷却速度の増加によってＲ2Ｔ14Ｂ結晶の粒径が減少し、チル晶と呼ばれる細かな等軸結晶の割合が増加するなどの問題があった。
本発明はＲ−リッチ相が細かく、かつ分散性がよく、またＲ2Ｔ14Ｂ結晶サイズが大きいＲ−Ｔ−Ｂ系焼結磁石用合金塊を提供することを目的とする。
Although the dispersibility of the R-rich phase is improved by the method of Patent Document 6, the temperature of the deposited alloy lump rises when the molten metal droplets are deposited, and the R-rich phase aggregates to form a pool. Therefore, in the fine pulverization step of the sintered magnet manufacturing process, the R-rich phase is first pulverized, which causes a problem that the time variation of the obtained powder composition is not stable. In addition, the dispersibility of the R-rich phase of the obtained powder was worse than that of a cast piece (hereinafter abbreviated as SC alloy) cast by the SC method, and thus there was a problem that the coercive force was also low.
Further, the method of Patent Document 7 has a problem that the grain size of the R 2 T 14 B crystal decreases due to an increase in the cooling rate, and the proportion of fine equiaxed crystals called chill crystals increases.
An object of the present invention is to provide an alloy block for an RTB-based sintered magnet having a fine R-rich phase, good dispersibility, and a large R 2 T 14 B crystal size.
本発明者は、ＮＣＣ法について改良を重ね、鋳型内面状態と溶湯の供給速度を最適化することで、高保磁力で且つ高配向率を有し着磁特性も良好な焼結磁石として最適の組織の合金塊を発明した。すなわち本発明は、
（１）Ｒ2Ｔ14Ｂ柱状結晶とＲ−リッチ相を含むＲ−Ｔ−Ｂ系希土類磁石用合金（ＲはＹを含む希土類元素の少なくとも１種、ＴはＦｅまたはＦｅとＦｅ以外の遷移金属元素の少なくとも１種、Ｂはボロンまたはボロンと炭素）であって、鋳造のままの状態で、Ｒ−リッチ相が断面でほぼ線状ないし棒状（これらの幅方向が短軸方向）に分散し、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが５００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とするＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（２）本発明は、Ｒ2Ｔ14Ｂ柱状結晶とＲ−リッチ相を含むＲ−Ｔ−Ｂ系希土類磁石用合金（ＲはＹを含む希土類元素の少なくとも１種、ＴはＦｅまたはＦｅとＦｅ以外の遷移金属元素の少なくとも１種、Ｂはボロンまたはボロンと炭素）であって、鋳造のままの状態で、短軸方向の長さが５μｍ以上であるＲリッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが５００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とするＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
The inventor repeatedly improved the NCC method, and optimized the inner surface state of the mold and the supply rate of the molten metal, so that the optimum structure as a sintered magnet having a high coercive force, a high orientation ratio, and a good magnetization characteristic was obtained. Invented an alloy lump. That is, the present invention
(1) R-T-B system rare earth magnet alloy including R 2 T 14 B columnar crystals and R-rich phase (R is at least one rare earth element including Y, T is transition of Fe or transition other than Fe and Fe) At least one kind of metal element, B is boron or boron and carbon), and the R-rich phase is distributed in a substantially linear or rod-like cross section in the as-cast state (the width direction is the minor axis direction). In addition, the area ratio of the region in which the length in the major axis direction of the R 2 T 14 B columnar crystal grains is 500 μm or more and the length in the minor axis direction is 50 μm or more is 10% or more of the whole alloy. It is an alloy lump for R-T-B system sintered magnet.
(2) The present invention provides an R-T-B system rare earth magnet alloy containing R 2 T 14 B columnar crystals and an R-rich phase (R is at least one rare earth element containing Y, and T is Fe or Fe. At least one transition metal element other than Fe, B is boron or boron and carbon), and the area ratio of the R-rich phase with the length in the minor axis direction of 5 μm or more is as-cast in the alloy. 10% or less of all R-rich phases present in the region, and the length of the R 2 T 14 B columnar crystal grains in the major axis direction is 500 μm or more and the length in the minor axis direction is 50 μm or more. Is an alloy ingot for an R-T-B system sintered magnet characterized by having an area ratio of 10% or more of the whole alloy.
（３）本発明は、短軸方向の長さが５μｍ以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが１０００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とする上記（１）または（２）に記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（４）本発明は、短軸方向の長さが５μｍ以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが１０００μｍ以上であり短軸方向の長さが１００μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とする上記（１）乃至（３）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（５）本発明は、短軸方向の長さが３μｍ以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが５００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とする上記（１）または（２）に記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（６）本発明は、短軸方向の長さが３μｍ以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが１０００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とする上記（１）乃至（３）のいずれか、または上記（５）に記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
(3) In the present invention, the area ratio of the R-rich phase having a minor axis length of 5 μm or more is 10% or less of all R-rich phases present in the alloy, and R 2 T 14 (1) or (1) above, wherein the area ratio of the B columnar crystal grains having a major axis length of 1000 μm or more and a minor axis direction length of 50 μm or more is 10% or more of the whole alloy. It is an alloy lump for RTB-based sintered magnets as described in 2).
(4) In the present invention, the area ratio of the R-rich phase whose minor axis length is 5 μm or more is 10% or less of all the R-rich phases present in the alloy, and R 2 T 14 (1) to (1) above, wherein the area ratio of the region where the length of the major axis direction of the B columnar crystal grains is 1000 μm or more and the length of the minor axis direction is 100 μm or more is 10% or more of the whole alloy. It is an alloy lump for RTB-based sintered magnets according to any one of 3).
(5) In the present invention, the area ratio of the R-rich phase having a minor axis length of 3 μm or more is 10% or less of all the R-rich phases present in the alloy, and R 2 T 14 (1) or (1) above, wherein the area ratio of the region where the length of the major axis direction of the B columnar crystal grains is 500 μm or more and the length of the minor axis direction is 50 μm or more is 10% or more of the whole alloy It is an alloy lump for RTB-based sintered magnets as described in 2).
(6) In the present invention, the area ratio of the R-rich phase having a minor axis length of 3 μm or more is 10% or less of all the R-rich phases present in the alloy, and R 2 T 14 (1) to (1) above, wherein the area ratio of the B columnar crystal grains having a length in the major axis direction of 1000 μm or more and a length in the minor axis direction of 50 μm or more is 10% or more of the whole alloy. 3), or an alloy block for an RTB-based sintered magnet according to (5) above.
（７）本発明は、短軸方向の長さが３μｍ以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であり、且つ、Ｒ2Ｔ14Ｂ柱状結晶粒の長軸方向の長さが１０００μｍ以上であり短軸方向の長さが１００μｍ以上である領域の面積率が合金全体の１０％以上あることを特徴とする上記（１）乃至（６）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（８）本発明は、Ｒ−リッチ相の間隔が平均１０μｍ以下であることを特徴とする上記（１）乃至（７）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（９）本発明は、Ｒ−リッチ相のアスペクト比が１０以上である上記（１）乃至（８）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（１０）本発明は、Ｒリッチ相の長さが平均５０〜１００μｍである上記（１）乃至（９）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（１１）本発明は、α−Ｆｅが実質的に無いことを特徴とする上記（１）乃至（１０）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
(7) In the present invention, the area ratio of the R-rich phase having a minor axis length of 3 μm or more is 10% or less of all the R-rich phases present in the alloy, and R 2 T 14 (1) to (1) above, wherein the area ratio of the region where the length of the major axis direction of the B columnar crystal grains is 1000 μm or more and the length of the minor axis direction is 100 μm or more is 10% or more of the whole alloy. 6) The alloy block for an RTB-based sintered magnet according to any one of 6).
(8) The present invention relates to the alloy block for an R-T-B system sintered magnet according to any one of (1) to (7) above, wherein an interval between R-rich phases is 10 μm or less on average. It is.
(9) The present invention is the alloy block for an RTB-based sintered magnet according to any one of the above (1) to (8), wherein the aspect ratio of the R-rich phase is 10 or more.
(10) The present invention is the alloy block for an RTB-based sintered magnet according to any one of (1) to (9), wherein the length of the R-rich phase is 50 to 100 μm on average.
(11) The present invention provides the alloy block for an RTB-based sintered magnet according to any one of the above (1) to (10), which is substantially free of α-Fe.
（１２）本発明は、厚さが１ｍｍ以上である上記（１）乃至（１１）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊である。
（１３）本発明は、溶湯を回転体に注ぎ、該回転体の回転によって溶湯を飛散させ、その飛散した溶湯を円筒状鋳型の内面で堆積凝固させる遠心鋳造法により鋳造する上記（１）乃至（１２）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊の製造法である。
（１４）本発明は、上記（１３）に記載の合金塊の製造法において、回転体の回転軸Ｒと円筒状鋳型の回転軸Ｌとが平行でない遠心鋳造法である上記（１）乃至（１２）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊の製造法である。
(12) The present invention is the alloy block for an RTB-based sintered magnet according to any one of (1) to (11), wherein the thickness is 1 mm or more.
(13) According to the present invention, the molten metal is poured into a rotating body, the molten metal is scattered by the rotation of the rotating body, and the scattered molten metal is cast by a centrifugal casting method that deposits and solidifies on the inner surface of the cylindrical mold. It is a manufacturing method of the alloy lump for RTB system sintered magnets in any one of (12).
(14) The present invention relates to the above-mentioned (1) to (1), wherein the alloy ingot manufacturing method according to (13) is a centrifugal casting method in which the rotation axis R of the rotating body and the rotation axis L of the cylindrical mold are not parallel. 12) The method for producing an alloy block for an RTB-based sintered magnet according to any one of 12).
（１５）本発明は、上記（１３）または（１４）に記載の合金塊の製造法において、円筒状鋳型が、その内壁面に該鋳型の材質よりも熱伝導率が小さい膜が設けられている遠心鋳造法である上記（１）乃至（１２）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊の製造法である。
（１６）本発明は、上記（１３）乃至（１５）のいずれかに記載の合金塊の製造法において、鋳造の初めに鋳造速度を早くし、その後遅くすることを特徴とする上記（１）乃至（１２）のいずれかに記載のＲ−Ｔ−Ｂ系焼結磁石用合金塊の製造法である。
（１７）本発明は、上記（１）乃至（１２）のいずれかに記載の合金塊を原料として製造したＲ−Ｔ−Ｂ系焼結磁石である。
(15) The present invention provides the method for producing an alloy ingot described in (13) or (14) above, wherein the cylindrical mold is provided with a film having a lower thermal conductivity on the inner wall surface than the material of the mold. A method for producing an alloy block for an RTB-based sintered magnet according to any one of the above (1) to (12), which is a centrifugal casting method.
(16) In the method for producing an alloy ingot according to any one of (13) to (15), the present invention is characterized in that the casting speed is increased at the beginning of casting and then is decreased thereafter. It is a manufacturing method of the alloy lump for RTB system sintered magnets in any one of thru | or (12).
(17) The present invention is an RTB-based sintered magnet manufactured using the alloy ingot according to any one of (1) to (12) above as a raw material.
本発明の合金塊は従来見られないＲ−リッチ相の細かさと均一性と柱状晶が大きいことを併せ持ち、本合金塊から製造した焼結磁石は従来より高特性、即ち高保磁力、高配向率及び良好な着磁特性を発現する。 The alloy lump of the present invention has the fineness and uniformity of the R-rich phase, which is not seen in the past, and the large number of columnar crystals, and the sintered magnet produced from this alloy lump has higher characteristics than before, that is, high coercive force and high orientation rate. And good magnetizing properties.
例えばＮｄ−Ｆｅ−Ｂ系のＳＣ合金（Ｎｄ３２．０質量％）の断面をＳＥＭ（走査電子顕微鏡）にて観察した時の反射電子像を図１に示す。図１の左側の面がロール面であり、右側の面が自由面である。ロール面から自由面までの長さ、すなわち鋳片の厚みは０．３ｍｍである。
白い部分が、Ｎｄ−リッチ相で（ＲがＮｄになっているためＲ−リッチ相をＮｄ−リッチ相と呼ぶ。）、その形状は凝固方向（左：ロール面側から右：自由面側）に向かって棒状に一部は繋がって延びているものと点状に散在しているものがある。棒状のものの長手方向は、ほぼ結晶の粒界や粒内でも結晶の成長方向に伸びている。Ｎｄ−リッチ相の融点は組成によって変化するが一般に６５０〜７５０℃と低い。そのためＮｄ2Ｆｅ14Ｂ相凝固後も、液相として存在し、冷却工程で若干消失あるいは分断されるものの、鋳造時の影響がそのままの形態で残っており、点状、線状、棒状のものが不均一に分布している。これは、ＳＣ法にて鋳造したＲ−Ｔ−Ｂ系合金塊の一般的な断面組織を表している。
なお、図１で線状、棒状に見えるＮｄ−リッチ相は実際には板状（ラメラ状）であり、図１では板状のＮｄ−リッチ相をある方向で切り取った面を見ているため、線状、棒状に見える。
For example, FIG. 1 shows a reflected electron image when a cross section of an Nd—Fe—B-based SC alloy (Nd 32.0 mass%) is observed with an SEM (scanning electron microscope). The left surface in FIG. 1 is a roll surface, and the right surface is a free surface. The length from the roll surface to the free surface, that is, the thickness of the slab is 0.3 mm.
The white part is the Nd-rich phase (the R-rich phase is called the Nd-rich phase because R is Nd), and the shape is the solidification direction (left: roll surface side to right: free surface side) Some of them are connected in a rod shape toward the surface and others are scattered in a dot shape. The longitudinal direction of the rod-like material extends in the crystal growth direction almost at the grain boundaries and within the grains. The melting point of the Nd-rich phase varies depending on the composition but is generally as low as 650 to 750 ° C. Therefore, even after solidification of the Nd 2 Fe 14 B phase, it remains as a liquid phase and disappears or breaks down slightly in the cooling process, but the effect during casting remains as it is, and it is in the form of dots, lines, rods Are unevenly distributed. This represents a general cross-sectional structure of an RTB-based alloy ingot cast by the SC method.
Note that the Nd-rich phase that appears to be linear or rod-like in FIG. 1 is actually plate-like (lamellar), and in FIG. 1, the surface obtained by cutting the plate-like Nd-rich phase in a certain direction is seen. Looks like a line or rod.
また上述のＳＣ合金の断面を、磁気カー効果を利用した偏光顕微鏡で撮った写真を図２に示す。写真の左側の面がロール面であり、右側の面が自由面である。
ロール面近傍の一部にはチル晶と呼ばれる数μｍ程度の大きさのＮｄ2Ｆｅ14Ｂ等軸結晶（以下、等軸晶と称す）部分が見られるものの、概略、ロール面側から自由面側への凝固方向に向ってＮｄ2Ｆｅ14Ｂ柱状結晶（以下、柱状晶と称す）が伸びている。これはＲ−Ｔ−Ｂ系ＳＣ合金で一般的に見られ、柱状晶の短軸方向の長さの平均は１５〜２５μｍである。
FIG. 2 shows a photograph of a cross section of the above SC alloy taken with a polarizing microscope using the magnetic Kerr effect. The left side of the photo is the roll side, and the right side is the free side.
Although a portion near the roll surface has an Nd 2 Fe 14 B equiaxed crystal (hereinafter, referred to as equiaxed crystal) portion called chill crystal having a size of about several μm, it is roughly free from the roll surface side. Nd 2 Fe 14 B columnar crystals (hereinafter referred to as columnar crystals) extend in the direction of solidification toward the side. This is generally seen in the R-T-B system SC alloy, and the average length of the columnar crystals in the minor axis direction is 15 to 25 μm.
本発明の合金塊はＲ−Ｔ−Ｂ系（ＲはＹを含む希土類元素の少なくとも１種、ＴはＦｅまたはＦｅとＦｅ以外の遷移金属元素の少なくとも１種、Ｂはボロンまたはボロンと炭素）で、一般的にはＲが２８〜３５質量％、Ｂが０．８〜１．３質量％、残部がＴである。
本発明の合金塊（Ｎｄ３２．０質量％）の断面をＳＥＭにて観察した時の反射電子写真を図３に示す。図３の倍率は図１の倍率と同じである。図１と同様に、線状、棒状のＮｄ−リッチ相が図３の左側から右側に伸びている。
本発明の合金塊における第１の特徴は、図３に示すように、Ｒ−リッチ相がほとんど線状、棒状に均一に分散していることである。この線状、棒状のアスペクト比（長軸方向の長さ／短軸方向の長さ）が１０以上、好ましくは１５以上、より好ましくは２０以上、さらに好ましくは２５以上であるＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以上、好ましくは３０％以上あることである。全てのＲ−リッチ相は合金中での面積率は、合金の組成によって変化し最大でも３０％程度であり、最小は１％程度である。上記のＲ−リッチ相により、微粉砕時の粉体組成の時間変動が安定し、また粉体のＲ−リッチ相の分散性もＳＣ合金と同程度に良くなり、このため焼結性が向上して保磁力が高くなる。
The alloy ingot of the present invention is an R-T-B system (R is at least one rare earth element including Y, T is at least one of transition metal elements other than Fe or Fe and Fe, and B is boron or boron and carbon). In general, R is 28 to 35% by mass, B is 0.8 to 1.3% by mass, and the balance is T.
FIG. 3 shows a reflection electron photograph when a cross section of the alloy lump (Nd 32.0 mass%) of the present invention is observed with an SEM. The magnification in FIG. 3 is the same as that in FIG. Similar to FIG. 1, linear and rod-like Nd-rich phases extend from the left side to the right side in FIG.
The first feature of the ingot of the present invention is that, as shown in FIG. 3, the R-rich phase is almost uniformly distributed in a linear or rod shape. The R-rich phase having a linear or rod-like aspect ratio (length in the major axis direction / length in the minor axis direction) of 10 or more, preferably 15 or more, more preferably 20 or more, and even more preferably 25 or more. The area ratio is 10% or more, preferably 30% or more of all R-rich phases present in the alloy. The area ratio of all R-rich phases in the alloy varies depending on the composition of the alloy and is about 30% at the maximum, and about 1% at the minimum. The above R-rich phase stabilizes the time variation of the powder composition during fine pulverization, and the dispersibility of the R-rich phase of the powder is improved to the same extent as the SC alloy, which improves the sinterability. As a result, the coercive force is increased.
なお、図３で線状、棒状に見えるＮｄ−リッチ相は実際には板状（ラメラ状）であり、写真では板状のＮｄ−リッチ相をある方向で切り取って見ているため、線状、棒状に見える。 Note that the Nd-rich phase that appears to be linear or rod-like in FIG. 3 is actually plate-like (lamellar), and in the photograph, the plate-like Nd-rich phase is cut out in a certain direction, so that it is linear. Looks like a stick.
本発明の合金塊における特徴を別の視点から見れば、合金塊がＲ−リッチ相の融点以上の温度である程度の時間晒された場合に見られる、線状、棒状のＲ−リッチ相が凝集して短軸方向の長さが５μｍ以上に太くなるものもあるが、その場合でもＲ−リッチ相の短軸方向の長さが５μｍ以上のものの面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であることである。より好ましくは短軸方向の長さが３μｍ以上まで太くなったＲ−リッチ相の面積率が合金中に存在する全てのＲ−リッチ相の１０％以下であることである。そしてアスペクト比は上記した範囲のものが好ましい。 From another viewpoint, the characteristics of the alloy ingot of the present invention are agglomerated linear and rod-like R-rich phases, which are observed when the alloy ingot is exposed to a temperature equal to or higher than the melting point of the R-rich phase for a certain period of time. Although the length in the minor axis direction may be thicker than 5 μm, even in that case, the area ratio of the R-rich phase with the minor axis length in the minor axis direction of 5 μm or more is all R— present in the alloy. It is 10% or less of the rich phase. More preferably, the area ratio of the R-rich phase whose length in the minor axis direction is increased to 3 μm or more is 10% or less of all the R-rich phases present in the alloy. The aspect ratio is preferably within the above range.
また、本発明の合金塊は図３に示すように、Ｒ−リッチ相が層状におおよそ５０〜１００μｍごとに切れている部分が目立つことも特徴の一つである。これは後述する製造法に因るものであり、溶湯が厚み約５０〜１００μｍの板状になって堆積するためである。 In addition, as shown in FIG. 3, the alloy lump of the present invention is also characterized in that the portion where the R-rich phase is cut in layers is cut approximately every 50 to 100 μm. This is because the molten metal is deposited in the form of a plate having a thickness of about 50 to 100 μm.
ここでＲ−リッチ相の短軸方向の長さと面積率は例えば以下のようにして測定する。
合金塊断面を研磨し、ＳＥＭの反射電子像にて断面の任意の視野を４００倍にて、ランダムに１０視野で写真を撮る。それぞれの写真について画像処理を行い、各Ｒ−リッチ相の面積と、図５に示すように短軸方向の長さが３μｍ以上または５μｍ以上ある部分の面積を求める。ここで図５における任意の点Ｐでの短軸方向の長さは、点Ｐから図５のような線を引き、最も短い線の長さを単軸方向の長さとする（図５では実線）。
そして１０視野全てのＲ−リッチ相の面積を合計し、また、短軸方向の長さが３μｍ以上または５μｍ以上ある部分のＲ−リッチ相の面積を合計し、これらの数値の比を面積率とする。
または、写真をコピーし、このコピーした紙を切り取ってそれぞれの部分の重量を計る方法で求めてもよい。
なお、図６のような枝分かれしているように見える形状のＲ−リッチ相については、枝分かれの根元（点線の位置）で分割してそれぞれ別のＲ−リッチ相として画像処理する。
Here, the length and the area ratio of the R-rich phase in the minor axis direction are measured as follows, for example.
The alloy lump cross section is polished, and an arbitrary field of view of the cross section is taken 400 times in the backscattered electron image of the SEM, and photographs are taken at 10 fields at random. Image processing is performed on each photograph, and the area of each R-rich phase and the area of the portion having a length in the minor axis direction of 3 μm or more or 5 μm or more are obtained as shown in FIG. Here, the length in the minor axis direction at an arbitrary point P in FIG. 5 is drawn from the point P as shown in FIG. 5, and the length of the shortest line is the length in the uniaxial direction (the solid line in FIG. 5). ).
The total area of the R-rich phase of all 10 fields of view is summed, and the area of the R-rich phase of the portion having a length in the minor axis direction of 3 μm or more or 5 μm or more is summed. And
Or you may obtain | require by the method of copying a photograph, cutting out this copied paper, and measuring the weight of each part.
Note that the R-rich phase that looks like a branch as shown in FIG. 6 is divided at the root of the branch (the position of the dotted line) and subjected to image processing as a separate R-rich phase.
本発明の合金塊の断面を、磁気カー効果を利用した偏光顕微鏡で撮った写真を図４に示す。図４の倍率は図２の倍率と同じである。柱状晶は厚み方向にほぼ沿って伸びており、図４はその一部を撮ったものである。
本発明の合金塊における第２の特徴は、それぞれの柱状晶の面積が、図２に示すＳＣ合金の柱状晶の面積よりも大きいことである。つまり、長軸方向の長さが５００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上、好ましくは３０％以上あることである。また好ましくは、長軸方向の長さが１０００μｍ以上であり短軸方向の長さが５０μｍ以上である領域の面積率が合金全体の１０％以上好ましくは２０％以上あることである。さらに好ましくは、長軸方向の長さが１０００μｍ以上であり短軸方向の長さが１００μｍ以上である領域の面積率が合金全体の１０％以上好ましくは２０％以上あることである。このことにより、微粉砕工程で得られる結晶方位が１方向しかない粉体が増えるため、焼結磁石にした場合の配向率が高くなる。
FIG. 4 shows a photograph of a cross section of the alloy lump of the present invention taken with a polarization microscope utilizing the magnetic Kerr effect. The magnification of FIG. 4 is the same as that of FIG. The columnar crystal extends substantially along the thickness direction, and FIG. 4 is a part of it.
The second feature of the alloy ingot of the present invention is that the area of each columnar crystal is larger than the area of the columnar crystal of the SC alloy shown in FIG. That is, the area ratio of the region in which the length in the major axis direction is 500 μm or more and the length in the minor axis direction is 50 μm or more is 10% or more, preferably 30% or more of the entire alloy. Also preferably, the area ratio of the region having a length in the major axis direction of 1000 μm or more and a length in the minor axis direction of 50 μm or more is 10% or more, preferably 20% or more of the whole alloy. More preferably, the area ratio of the region where the length in the major axis direction is 1000 μm or more and the length in the minor axis direction is 100 μm or more is 10% or more, preferably 20% or more of the entire alloy. This increases the amount of powder obtained in the fine pulverization process in which the crystal orientation is only in one direction, so that the orientation ratio in the case of a sintered magnet is increased.
ここで結晶粒の長軸方向の長さ、短軸方向の長さ、面積率は例えば以下のようにして測定する。
合金塊断面を研磨し、偏光顕微鏡にて断面の任意の３箇所で合金の一端から他端まで厚み方向に沿って５０倍で連続写真を撮る。それぞれの連続写真において、柱状晶の長軸方向の長さが５００μｍ以上、または１０００μｍ以上ある柱状晶を特定する。さらにこれらの柱状晶において、短軸方向の長さが５０μｍ、または１００μｍ以上ある部分の面積を求める。３箇所の連続写真について求めたこれらの面積を、３箇所の連続写真の全断面積の合計で除することで所定の面積率を求めることができる。
なお、それぞれの面積は画像処理で求めてもよいし、写真をコピーし、このコピーした紙を切り取ってそれぞれの部分の重量を計る方法で求めてもよい。
Here, the length in the major axis direction, the length in the minor axis direction, and the area ratio of the crystal grains are measured as follows, for example.
The alloy lump cross section is polished, and a continuous photograph is taken at a magnification of 50 along the thickness direction from one end of the alloy to the other end at any three locations in the cross section with a polarizing microscope. In each continuous photograph, a columnar crystal having a length in the major axis direction of the columnar crystal of 500 μm or more, or 1000 μm or more is specified. Further, in these columnar crystals, the area of the portion having a length in the minor axis direction of 50 μm or 100 μm or more is obtained. A predetermined area ratio can be obtained by dividing these areas obtained for three consecutive photographs by the sum of the total cross-sectional areas of the three consecutive photographs.
Each area may be obtained by image processing, or may be obtained by copying a photograph, cutting out the copied paper, and measuring the weight of each part.
本発明の合金塊における第３の特徴は、Ｒ−リッチ相間隔が平均１０μｍ以下であることである。このことにより第１の特徴に加えて微粉砕後のＲ−リッチ相の分散性が向上し、焼結性が向上することで保磁力が向上する。
Ｒ−リッチ相間隔は、合金塊の断面をＳＥＭにて観察し、鋳造厚さ方向と直角方向のＲ−リッチ相間隔を、画像処理あるいは写真上からの手測定により平均したものである。
The third feature of the alloy ingot of the present invention is that the R-rich phase interval is 10 μm or less on average. Thereby, in addition to the first feature, the dispersibility of the R-rich phase after pulverization is improved, and the coercivity is improved by improving the sinterability.
The R-rich phase interval is obtained by observing the cross section of the alloy lump with an SEM, and averaging the R-rich phase interval in the direction perpendicular to the casting thickness direction by image processing or manual measurement from a photograph.
さらに、本発明の合金塊における第４の特徴は、Ｒ成分が化学量論組成付近まで実質的にα―Ｆｅが発生しないことである。ここで実質的にα―Ｆｅが発生しないというのは、合金塊の任意の断面の任意の視野で１０視野にてα―Ｆｅが存在するかどうかを確認した場合９割以上の視野で見つからない程度の状態をいう。ＳＥＭの反射電子線像では、α―Ｆｅはデンドライト状に黒く見える。 Furthermore, the fourth feature of the alloy ingot of the present invention is that α-Fe is not substantially generated until the R component is close to the stoichiometric composition. Here, the fact that α-Fe does not substantially occur is not found in 90% or more of visual fields when it is confirmed whether or not α-Fe is present in 10 visual fields of arbitrary cross sections of the alloy ingot. The state of the degree. In the reflected electron beam image of SEM, α-Fe looks black like a dendrite.
本発明の合金塊は、次のような方法にて製造できる。図７は、本発明の１例であり、これを用いて説明する。 The alloy lump of the present invention can be manufactured by the following method. FIG. 7 shows an example of the present invention, which will be described below.
通常、希土類合金は、その活性な性質なため真空または不活性ガスの部屋１の中でルツボ３にて溶解される。溶湯３１は、湯道６により回転軸をＲとした回転体５に受け、該回転体５の回転によって溶湯を円筒状の鋳型４の内壁に飛散させる。回転体は、回転軸をＲとして回転する物質であり、注がれた溶湯を周囲に飛散させる機能を有する物体であり、円盤、上に角度を持つカップ状、下に角度を持つコーン状等にて飛散させられるが、図で示すような容器状で側面に複数の孔部１１を有する形状（回転受け容器）が好ましい。 Normally, rare earth alloys are melted in a crucible 3 in a vacuum or inert gas chamber 1 due to their active nature. The molten metal 31 is received by the rotating body 5 having a rotation axis R by the runner 6, and the molten body is scattered on the inner wall of the cylindrical mold 4 by the rotation of the rotating body 5. The rotating body is a substance that rotates with the rotation axis as R, and is an object having a function of scattering the poured molten metal around, such as a disk, a cup shape with an angle on top, a cone shape with an angle on the bottom, etc. However, a shape having a plurality of holes 11 on the side surface (rotary receiving container) as shown in the figure is preferable.
このような回転体や回転体の内部に溶湯が注がれた場合、溶湯は回転による力や遠心力により、回転体の周囲に飛散させられる。この場合、回転体の熱容量を小さくするか十分余熱することによって溶湯を回転体上で凝固させず、円筒上鋳型の内壁にて堆積凝固させることができる。
図７では、鋳型が水平に置かれているが、垂直に置いても、傾斜させておいても回転体との位置関係を一定に保てばなんら問題はない。
When the molten metal is poured into such a rotating body or the rotating body, the molten metal is scattered around the rotating body by the force of rotation or centrifugal force. In this case, the molten metal can be deposited and solidified on the inner wall of the mold on the cylinder without reducing the heat capacity of the rotating body or sufficiently heating it, without solidifying the molten metal on the rotating body.
In FIG. 7, the mold is placed horizontally, but there is no problem if the positional relationship with the rotating body is kept constant regardless of whether it is placed vertically or inclined.
回転体５の回転軸Ｒと鋳型４の回転軸Ｌは、平行にすることも可能であるが、ある角度θをもたせることにより堆積面を鋳型の長手方向全体に広げることができ、それによって溶湯の堆積速度をコントロールすることが出来る。
この角度をつけることにより、溶湯を大きな面積範囲にばら撒くことができ結果的に凝固速度を大きくすることができる。
溶湯を鋳型内全体にばら撒くには、上述の角度をつける方法以外に、鋳型又は回転体を鋳型回転軸方法に前後させることによっても同様の効果が得られる。
The rotation axis R of the rotating body 5 and the rotation axis L of the mold 4 can be made parallel, but by having an angle θ, the deposition surface can be spread over the entire longitudinal direction of the mold, thereby The deposition rate can be controlled.
By providing this angle, the molten metal can be spread over a large area range, and as a result, the solidification rate can be increased.
In order to spread the molten metal throughout the mold, the same effect can be obtained by moving the mold or the rotating body back and forth to the mold rotating shaft method in addition to the above-described method of setting the angle.
なお、回転体と鋳型は同一方向に回転速度をずらして回転させることが好ましい。反対方向に回転させると、溶湯が鋳型に衝突する際に鋳型に乗らずに飛散するスプラッシュ現象が発生し易くなり、歩留りの低下を招く。
また、回転体と鋳型の回転速度が同じであると鋳型上の同一面に線状に堆積することになり、鋳型前面に広がらない。
従ってあまり両者の回転速度が近いことも避けるべきで、通常は、両者の回転速度の差は少なくとも１０％以上、望ましくは２０％以上差をつけるべきである。
Note that the rotating body and the mold are preferably rotated at different rotational speeds in the same direction. When it is rotated in the opposite direction, when the molten metal collides with the mold, a splash phenomenon is likely to occur without getting on the mold and the yield is reduced.
Further, if the rotational speeds of the rotating body and the mold are the same, they will be deposited linearly on the same surface on the mold and will not spread on the front surface of the mold.
Therefore, it should be avoided that the rotational speeds of the two are too close. Normally, the difference between the rotational speeds of the two should be at least 10% or more, preferably 20% or more.
回転体の回転数は、溶湯の遠心力により溶湯が鋳型の内壁面に衝突するような条件を選ぶ必要がある。また、鋳型の回転数は、堆積凝固した合金塊が落下しないように１Ｇ以上の遠心力を与えるとともに、遠心力を増すことにより溶湯を鋳型内壁へ押し付けることで冷却効果を増すことができる。 The rotational speed of the rotating body needs to be selected so that the molten metal collides with the inner wall surface of the mold due to the centrifugal force of the molten metal. Further, the rotational speed of the mold can increase the cooling effect by applying a centrifugal force of 1 G or more so that the deposited and solidified alloy lump does not fall and pressing the molten metal against the inner wall of the mold by increasing the centrifugal force.
本発明の特徴は、鋳型内面に衝突させた溶湯を直ぐに凝固させるのではなく、一時的に液相線温度以上に保持することで、先に堆積した合金の結晶方位に沿って結晶化させ、その後堆積し一体化した合金の温度をＲ−リッチ相の融点をあまり越えない温度以下に保つことである。液相線温度は溶湯のＲ成分によって違うがおおよそ１１５０〜１３００℃である。この液相線温度以上に保持する時間は、好ましくは０．００１〜１秒である。さらに好ましくは０．００１〜０．１秒である。このことにより、γ―Ｆｅを発生させることなく短軸方向に長さが大きい柱状晶を成長させることができる。またＲ−リッチ相の融点もＲ成分によって違うが、おおよそ６５０〜７５０℃であり、Ｒ−リッチ相の融点をあまり越えない温度とは融点より精々１００℃高い温度である。この温度を越えると、Ｒ−リッチ相が凝集し短軸方向の長さが太くなるとともに、Ｒ−リッチ相の分散性も悪くなる。 The feature of the present invention is that the molten metal collided with the inner surface of the mold is not immediately solidified, but temporarily maintained at a liquidus temperature or higher, so that it is crystallized along the crystal orientation of the previously deposited alloy, The temperature of the subsequently deposited and integrated alloy is kept below the temperature that does not greatly exceed the melting point of the R-rich phase. The liquidus temperature is approximately 1150 to 1300 ° C. although it varies depending on the R component of the molten metal. The time for maintaining the liquidus temperature or higher is preferably 0.001 to 1 second. More preferably, it is 0.001-0.1 second. This makes it possible to grow columnar crystals having a long length in the minor axis direction without generating γ-Fe. Also, the melting point of the R-rich phase varies depending on the R component, but is approximately 650 to 750 ° C. The temperature not exceeding the melting point of the R-rich phase is a temperature that is exactly 100 ° C. higher than the melting point. When this temperature is exceeded, the R-rich phase aggregates and the length in the short axis direction becomes thick, and the dispersibility of the R-rich phase also deteriorates.
なお、図３ではＲ−リッチ相がおよそ５０〜１００μｍ間隔で層状に切れているのに対し、図４では柱状晶はそのような層状に切れてはいない。上述の本発明の方法により、柱状晶を切れることなく成長させることができる。 In FIG. 3, the R-rich phase is cut into layers at intervals of about 50 to 100 μm, whereas in FIG. 4, the columnar crystals are not cut into such layers. By the method of the present invention described above, the columnar crystals can be grown without breaking.
通常１３００〜１５００℃の溶湯を、鋳型内面に衝突した時から堆積終了（鋳造終了）まででこのような温度変化をさせるには、鋳型内面と合金との間の熱伝達係数をできるだけ大きくする必要がある。そのための方法としては例えば、鋳型内面に鋳型材質よりも熱伝導率が低い材料で作った膜を付ける方法がある。膜の材質は金属でもセラミックスでもそれらの複合材でもよい。また膜の厚みは１μｍから１ｍｍが好ましい。より好ましくは１μｍから５００μｍである。堆積開始（鋳造開始）から数十秒間に多量の溶湯を堆積させることにより、合金の鋳型面の平滑度が向上し、熱伝達係数を大きくすることができる。別の見方をすれば、鋳型内面に熱伝導率の悪い膜を付けて熱伝導率を低くすることで、最初に堆積した合金塊の温度の冷却を悪くして、この合金塊に高温変形能があるうちに鋳型の遠心力により合金塊を鋳型に密着させて鋳型と合金塊との間の熱伝達係数を大きくする。この時、高温に保って変形させやすくするため、堆積速度を速くする（溶湯供給量を増やす）。その後は堆積速度を遅くする（溶湯供給量を減らす）ことで、鋳型への十分な熱移動時間を与えて合金内部の温度上昇を抑える。合金の厚みが増すほど鋳型への熱移動は遅くなるため、合金が厚くなるのに応じて堆積速度を遅くすることが好ましい。さらに好ましくは、最初の堆積の後の適切な短時間、後の堆積速度よりも遅くして最初に堆積した合金塊の熱を鋳型に移動させるのに十分な時間を与えることである。 In order to cause such a temperature change from when the molten metal of 1300 to 1500 ° C. collides with the inner surface of the mold until the end of deposition (casting), it is necessary to increase the heat transfer coefficient between the inner surface of the mold and the alloy as much as possible. There is. As a method therefor, for example, there is a method in which a film made of a material having a lower thermal conductivity than the mold material is attached to the inner surface of the mold. The material of the film may be metal, ceramics, or a composite material thereof. The thickness of the membrane is preferably 1 μm to 1 mm. More preferably, it is 1 μm to 500 μm. By depositing a large amount of molten metal within several tens of seconds from the start of deposition (start of casting), the smoothness of the mold surface of the alloy can be improved and the heat transfer coefficient can be increased. From another viewpoint, a low thermal conductivity film is attached to the inner surface of the mold to lower the thermal conductivity, so that the cooling of the temperature of the initially deposited alloy mass is worsened. In the meantime, the alloy lump is brought into close contact with the mold by the centrifugal force of the mold to increase the heat transfer coefficient between the mold and the alloy lump. At this time, the deposition rate is increased (the amount of molten metal supplied is increased) in order to keep the temperature high and to facilitate deformation. Thereafter, by slowing the deposition rate (decreasing the amount of molten metal supplied), sufficient heat transfer time to the mold is given to suppress the temperature rise inside the alloy. As the alloy thickness increases, the heat transfer to the mold becomes slower, so it is preferable to slow the deposition rate as the alloy becomes thicker. More preferably, an adequate short time after the initial deposition, slower than the subsequent deposition rate, to allow sufficient time to transfer the heat of the initially deposited alloy mass to the mold.
また、最初に堆積した合金塊の変形能を向上させ、またチル晶の生成を抑制するため、鋳型内面の温度を予め２００〜７５０℃で加熱しておいてもよい。２００℃より低いとこれらの効果は期待できない。一方、７５０℃を越えるとＲ−リッチ相の融点より高いことから、堆積した合金塊の温度が下がりにくくなり、Ｒ−リッチ相がプール化してしまう。 Moreover, in order to improve the deformability of the alloy lump deposited initially and to suppress the formation of chill crystals, the temperature of the inner surface of the mold may be preheated at 200 to 750 ° C. If the temperature is lower than 200 ° C., these effects cannot be expected. On the other hand, if it exceeds 750 ° C., it is higher than the melting point of the R-rich phase, so that the temperature of the deposited alloy lump is hardly lowered, and the R-rich phase pools.
鋳型の材質は、常温での熱伝導率が３０〜４１０Ｗｍ-1Ｋ-1の材質が好ましい。３０Ｗｍ-1Ｋ-1より低いと堆積した合金の冷却速度が遅くなり、Ｒ−リッチ相がプール化しやすくなる。一方、熱伝導率は大きいほど好ましいが、銀に代表される４１０Ｗｍ-1Ｋ-1を越える材料は高価であり、工業用には向かない。工業的には熱伝導率が大きい銅が好ましいが、鉄でも問題はない。
なお堆積開始時の堆積速度と堆積時間及び後の堆積速度は、溶湯の組成や鋳型材質、鋳型の回転軸方向、鋳型内面の遠心力、膜の熱伝導率などによって決まる最適値を選ぶ必要がある。
The material of the mold is preferably a material having a thermal conductivity of 30 to 410 Wm −1 K −1 at room temperature. If it is lower than 30 Wm −1 K −1 , the cooling rate of the deposited alloy becomes slow, and the R-rich phase tends to be pooled. On the other hand, the higher the thermal conductivity, the better, but a material exceeding 410 Wm −1 K −1 typified by silver is expensive and unsuitable for industrial use. Industrially, copper having a high thermal conductivity is preferable, but iron is no problem.
The deposition rate at the start of deposition, the deposition time, and the subsequent deposition rate should be selected to be optimal values determined by the molten metal composition, mold material, mold rotation axis direction, centrifugal force on the mold inner surface, thermal conductivity of the film, etc. is there.
合金の厚みは１ｍｍ以上あることが好ましい。１ｍｍ未満であまり薄すぎると生産性が悪くなる。 The thickness of the alloy is preferably 1 mm or more. If it is less than 1 mm and too thin, the productivity will deteriorate.
本鋳造法にて製造したＲ−Ｔ−Ｂ系磁石用合金塊から粉砕、成型、焼結することにより、高特性の異方性磁石を製造することができる。
粉砕は、通常、水素解砕、中粉砕、微粉砕の順で行なわれ、３μｍ（ＦＳＳＳ）程度の粉体にされる。
By pulverizing, molding, and sintering from an alloy block for an R-T-B magnet manufactured by the present casting method, an anisotropic magnet having high characteristics can be manufactured.
The pulverization is usually performed in the order of hydrogen pulverization, medium pulverization, and fine pulverization, and is made into a powder of about 3 μm (FSSS).
ここで、水素解砕は、前工程の水素吸蔵工程と後工程の脱水素工程に分けられる。水素吸蔵工程では、２０〜５０００ｋＰａの圧力の水素ガス雰囲気で、主に合金塊のＲ−リッチ相に水素を吸蔵させ、この時に生成されるＲ−水素化物によりＲ−リッチ相が体積膨張することを利用して、合金塊自体を微細に割ることまたは無数の微細な割れ目を生じさせる。この水素吸蔵は常温〜６００℃程度の範囲で実施されるが、Ｒ−リッチ相の体積膨張を大きくして効率良く割るためには、常温〜１００℃程度の範囲で実施することが好ましい。好ましい処理時間は１時間以上である。この水素吸蔵工程により生成したＲ−水素化物は大気中では不安定であり酸化され易いため、２００〜６００℃程度で１００Ｐａ以下真空中に保持する脱水素処理を行なうことが好ましい。この処理により、大気中で安定なＲ-水素化物に変化させることができる。好ましい処理時間は３０分以上である。水素吸蔵後から焼結までの各工程で酸化防止のための雰囲気管理がなされている場合は、脱水素処理を省くこともできる。
なお、この水素解砕をせずに中粉砕、微粉砕することもできる。
Here, hydrogen cracking is divided into a hydrogen storage process in the previous process and a dehydrogenation process in the subsequent process. In the hydrogen occlusion process, hydrogen is occluded mainly in the R-rich phase of the alloy lump in a hydrogen gas atmosphere at a pressure of 20 to 5000 kPa, and the R-rich phase is volume-expanded by the R-hydride generated at this time. Is used to finely break the alloy ingot itself or to generate countless fine cracks. This hydrogen occlusion is carried out in the range from room temperature to about 600 ° C., but it is preferably carried out in the range from room temperature to about 100 ° C. in order to increase the volume expansion of the R-rich phase and efficiently divide it. A preferred treatment time is 1 hour or more. Since the R-hydride produced by this hydrogen occlusion process is unstable in the atmosphere and easily oxidized, it is preferable to perform a dehydrogenation treatment in which the pressure is kept at about 200 to 600 ° C. in a vacuum of 100 Pa or less. By this treatment, it can be changed to R-hydride which is stable in the atmosphere. A preferred treatment time is 30 minutes or more. In the case where atmosphere management for preventing oxidation is performed in each process from hydrogen storage to sintering, dehydrogenation treatment can be omitted.
It is also possible to pulverize inside and finely without this hydrogen cracking.
中粉砕とは、合金片をアルゴンガスや窒素ガスなどの不活性ガス雰囲気中で、例えば５００μｍ以下まで粉砕することである。このための粉砕機には、例えばブラウンミル粉砕機がある。本発明の水素解砕した合金片の場合、既に微細に割れている、または内部に無数の微細な割れ目が生じているため、この中粉砕を省略することもできる。 Medium pulverization refers to pulverizing an alloy piece to, for example, 500 μm or less in an inert gas atmosphere such as argon gas or nitrogen gas. An example of a pulverizer for this purpose is a brown mill pulverizer. In the case of the hydrogen-crushed alloy piece of the present invention, since it has already been finely cracked or innumerable fine cracks are generated inside, the pulverization can be omitted.
微粉砕とは、３μｍ（ＦＳＳＳ）程度まで粉砕することである。このための粉砕機には、例えばジェットミル装置がある。この場合、粉砕時の雰囲気はアルゴンガスや窒素ガスなどの不活性ガス雰囲気とする。これらの不活性ガス中に２質量％以下、好ましくは１質量％以下の酸素を混入させてもよい。このことにより粉砕効率が向上するとともに、粉砕後の粉体の酸素濃度が１０００〜１００００ｐｐｍとなり耐酸化性が向上する。また、焼結時の異常粒成長を抑制することもできる。 The fine pulverization is to pulverize to about 3 μm (FSSS). An example of a pulverizer for this purpose is a jet mill device. In this case, the atmosphere during pulverization is an inert gas atmosphere such as argon gas or nitrogen gas. 2% by mass or less, preferably 1% by mass or less of oxygen may be mixed into these inert gases. This improves the pulverization efficiency, and the oxygen concentration of the pulverized powder becomes 1000 to 10,000 ppm, thereby improving the oxidation resistance. Also, abnormal grain growth during sintering can be suppressed.
磁場成型時に粉体と金型内壁との摩擦を低減し、また粉体どうしの摩擦も低減させて配向性を向上させるため、粉体にはステアリン酸亜鉛等の潤滑剤を添加することが好ましい。好ましい添加量は０．０１〜１質量％である。添加は微粉砕前でも後でもよいが、磁場中成形前に、アルゴンガスや窒素ガスなどの不活性ガス雰囲気中でＶ型ブレンダー等を用いて十分に混合することが好ましい。 It is preferable to add a lubricant such as zinc stearate to reduce the friction between the powder and the inner wall of the mold during magnetic field molding and to improve the orientation by reducing the friction between the powders. . A preferable addition amount is 0.01 to 1% by mass. The addition may be before or after fine pulverization, but it is preferable to sufficiently mix using a V-type blender or the like in an inert gas atmosphere such as argon gas or nitrogen gas before molding in a magnetic field.
３μｍ（ＦＳＳＳ）程度まで粉砕された粉体は、磁場中成型機でプレス成型される。金型は、キャビティ内の磁界方向を考慮して、磁性材と非磁性材を組み合わせて作製される。成型圧力は５０〜２００ＭＰａが好ましい。成型時のキャビティ内の磁界は４００〜１６００ｋＡｍ-1が好ましい。また、成型時の雰囲気はアルゴンガスや窒素ガスなどの不活性ガス雰囲気が好ましいが、上述の耐酸化処理した粉体の場合、大気中でも可能である。 The powder pulverized to about 3 μm (FSSS) is press-molded in a magnetic field molding machine. The mold is manufactured by combining a magnetic material and a non-magnetic material in consideration of the direction of the magnetic field in the cavity. The molding pressure is preferably 50 to 200 MPa. The magnetic field in the cavity at the time of molding is preferably 400-1600 kAm- 1 . Further, the atmosphere during molding is preferably an inert gas atmosphere such as argon gas or nitrogen gas, but in the case of the above-mentioned oxidation-resistant powder, it can also be performed in the air.
焼結は、１０００〜１１００℃で行なわれる。焼結温度に到達する前に潤滑剤と、微粉中の水素はできるだけ除去しておく必要がある。潤滑剤の好ましい除去条件は、１Ｐａ以下の真空中またはＡｒ減圧フロー雰囲気中、３００〜５００℃で３０分以上保持することである。また、水素の好ましい除去条件は、１Ｐａ以下の真空中、７００〜９００℃で３０分以上保持することである。焼結時の雰囲気はアルゴンガス雰囲気または１Ｐａ以下の真空雰囲気が好ましい。保持時間は１時間以上が好ましい。 Sintering is performed at 1000 to 1100 ° C. Before reaching the sintering temperature, it is necessary to remove the lubricant and hydrogen in the fine powder as much as possible. A preferable removal condition of the lubricant is to hold at 300 to 500 ° C. for 30 minutes or more in a vacuum of 1 Pa or less or in an Ar reduced pressure flow atmosphere. Moreover, the preferable removal conditions of hydrogen are hold | maintaining for 30 minutes or more at 700-900 degreeC in the vacuum of 1 Pa or less. The atmosphere during sintering is preferably an argon gas atmosphere or a vacuum atmosphere of 1 Pa or less. The holding time is preferably 1 hour or longer.
焼結後、保磁力向上のため、必要に応じて５００〜６５０℃で熱処理することができる。好ましい雰囲気はアルゴンガス雰囲気または真空雰囲気である。好ましい保持時間は３０分以上である。 After sintering, heat treatment can be performed at 500 to 650 ° C. as necessary to improve the coercive force. A preferable atmosphere is an argon gas atmosphere or a vacuum atmosphere. A preferable holding time is 30 minutes or more.
以下に実施例を挙げて説明するが、本発明は以下の実施例に限定されるものではない。
（実施例１）
合金組成が、Ｎｄ：２７．０質量％、Ｄｙ：５．０質量％、Ｂ：１．００質量％、Ｃｏ：１．０質量％、Ａｌ：０．３０質量％、Ｃｕ：０．１０質量％、残部鉄になるように、金属ネオジム、金属ジスプロジウム、フェロボロン、コバルト、アルミニウム、銅、鉄を配合し、アルミナ坩堝を使用して、アルゴンガス1気圧雰囲気中で、高周波溶解炉で溶解し、溶湯を図７に示す装置で鋳造を行った。
鋳型は内径５００ｍｍ、長さ５００ｍｍの鉄製とし、鋳型内面に８０Ｎｉ-２０Ｃｒの膜を１００μｍ溶射した。
回転受け容器は、直径２ｍｍの孔部を周囲に８個配置した内径２５０ｍｍのものである。回転受け容器の回転軸と鋳型の回転軸との角度θは２５度とした。
鋳型の回転数は、遠心力が３Ｇになるように、１０４ｒｐｍに設定し、回転受け容器の回転速度は５３５ｒｐｍとし、溶湯に約４０Ｇの遠心力を加えた。
鋳型内面への溶湯の平均堆積速度の条件は、堆積開始から１０秒間は０．３ｍｍ／秒、その後の１０秒間は０．２０ｍ／秒とし、その後終了までは０．１５ｍｍ／秒で一定とした。
得られた合金塊の厚さは、円筒状鋳型の中央部で８〜９ｍｍ、両端部近傍の最も厚い部分で１０〜１１ｍｍであった。また合金塊の鋳型面は平滑であった。
Examples will be described below, but the present invention is not limited to the following examples.
Example 1
Alloy composition is Nd: 27.0 mass%, Dy: 5.0 mass%, B: 1.00 mass%, Co: 1.0 mass%, Al: 0.30 mass%, Cu: 0.10 mass% %, And mixed with metal neodymium, metal dysprodium, ferroboron, cobalt, aluminum, copper, and iron so that the balance is iron. The molten metal was cast using the apparatus shown in FIG.
The mold was made of iron having an inner diameter of 500 mm and a length of 500 mm, and an 80Ni-20Cr film was sprayed on the inner surface of the mold by 100 μm.
The rotation receiving container has an inner diameter of 250 mm in which eight holes having a diameter of 2 mm are arranged around the periphery. The angle θ between the rotation axis of the rotating container and the rotation axis of the mold was set to 25 degrees.
The rotational speed of the mold was set to 104 rpm so that the centrifugal force was 3 G, the rotational speed of the rotating container was 535 rpm, and a centrifugal force of about 40 G was applied to the molten metal.
The condition of the average deposition rate of the melt on the inner surface of the mold was 0.3 mm / second for 10 seconds from the start of deposition, 0.20 m / second for the subsequent 10 seconds, and constant at 0.15 mm / second until the end. .
The thickness of the obtained alloy lump was 8 to 9 mm at the center of the cylindrical mold, and 10 to 11 mm at the thickest part near both ends. The mold surface of the alloy lump was smooth.
得られた合金塊のＲ−リッチ相については、ＳＥＭの反射電子像で任意の視野を４００倍にてランダムに１０視野で写真を撮り（図３がその一例である。図３で黒く見える部分は穴である。）、これら写真を画像処理して、短軸方向の長さが５μｍ以上、３μｍ以上であるＲ−リッチ相の面積率とＲ−リッチ相の平均間隔を測定した。
その結果、５μｍ以上の面積率は０％であり、３μｍ以上の面積率は４％であった。また、Ｒ−リッチ相の平均間隔は５μｍであった。
なおこの１０視野には、α―Ｆｅと思われる黒い相は無かった。
As for the R-rich phase of the obtained alloy lump, an SEM reflection electron image was taken at random 400 times and 10 pictures were taken at random (FIG. 3 is an example. The part that appears black in FIG. 3) ), These photographs were subjected to image processing, and the area ratio of the R-rich phase and the average interval between the R-rich phases having a length in the minor axis direction of 5 μm or more and 3 μm or more were measured.
As a result, the area ratio of 5 μm or more was 0%, and the area ratio of 3 μm or more was 4%. The average interval between the R-rich phases was 5 μm.
In these 10 fields of view, there was no black phase thought to be α-Fe.
また柱状晶については、偏光顕微鏡にて断面の任意の３箇所で合金の一端から他端まで厚み方向に沿って５０倍で連続写真を撮り（図４はその一部を拡大したものである。）、この写真を別紙にコピーした上で、このコピー紙を切り取って重さを計る手法で、柱状晶の長軸方向の長さが５００μｍ以上、または１０００μｍ以上あり、短軸方向の長さが５０μｍ、または１００μｍ以上ある部分の面積率を測定した。
その結果、長軸方向５００μｍ以上短軸方向５０μｍ以上は３８％、長軸方向１０００μｍ以上短軸方向１００μｍ以上は１６％であった。
As for the columnar crystals, continuous photographs were taken at a magnification of 50 times along the thickness direction from one end of the alloy to the other end at any three locations in the cross section with a polarizing microscope (FIG. 4 is an enlarged view of a part thereof. ) After copying this photograph to a separate sheet, the length of the long axis direction of the columnar crystal is 500 μm or more, or 1000 μm or more by cutting this copy paper and measuring the weight. The area ratio of the part which is 50 μm or 100 μm or more was measured.
As a result, the major axis direction was 500 μm or more and the minor axis direction 50 μm or more was 38%, and the major axis direction was 1000 μm or more and the minor axis direction 100 μm or more was 16%.
（比較例１）
実施例１と同様の組成の合金を配合し、実施例１と同様に溶解し、同様の鋳造装置にて鋳造を行った。
但し、鋳型内面には何らの膜も付けなかった。また、鋳型内面への溶湯の平均堆積速度の条件は、堆積開始から終了まで０．１５ｍｍ／秒で一定とした。
得られた合金塊の厚さは、円筒状鋳型の中央部で８〜９ｍｍ、両端部近傍の最も厚い部分で１０〜１１ｍｍであった。また合金塊の鋳型面は凹凸が酷く、深さ０.数ｍｍの穴が多数あった。
(Comparative Example 1)
An alloy having the same composition as in Example 1 was blended, dissolved in the same manner as in Example 1, and cast in the same casting apparatus.
However, no film was attached to the inner surface of the mold. The condition of the average deposition rate of the molten metal on the inner surface of the mold was constant at 0.15 mm / second from the start to the end of the deposition.
The thickness of the obtained alloy lump was 8 to 9 mm at the center of the cylindrical mold, and 10 to 11 mm at the thickest part near both ends. Also, the mold surface of the alloy lump was severely uneven, and had many holes with a depth of several millimeters.
得られた合金塊のＲ−リッチ相については、実施例１と同様の方法で短軸方向の長さが５μｍ以上、３μｍ以上であるＲ−リッチ相の面積率とＲ−リッチ相の平均間隔を測定した。
その結果、５μｍ以上の面積率は２２％であり、３μｍ以上の面積率は４１％であった。また、Ｒ−リッチ相の平均間隔は１３μｍであった。
なおこの１０視野には、α―Ｆｅと思われる黒い相は無かった。
About the R-rich phase of the obtained alloy lump, the area ratio of the R-rich phase whose average length in the minor axis direction is 5 μm or more and 3 μm or more and the average interval between the R-rich phases in the same manner as in Example 1. Was measured.
As a result, the area ratio of 5 μm or more was 22%, and the area ratio of 3 μm or more was 41%. The average interval between the R-rich phases was 13 μm.
In these 10 fields of view, there was no black phase thought to be α-Fe.
また柱状晶についても、実施例１と同様の方法で、柱状晶の長軸方向の長さ５００μｍ以上、または１０００μｍ以上あり、短軸方向の長さが５０μｍ、または１００μｍ以上ある部分の面積率を測定した。
その結果、長軸方向５００μｍ以上短軸方向５０μｍ以上は７２％、長軸方向１０００μｍ以上短軸方向１００μｍ以上は６８％であった。
For columnar crystals, the area ratio of the columnar crystal having a length in the major axis direction of 500 μm or more or 1000 μm or more and a length in the minor axis direction of 50 μm or 100 μm or more is determined in the same manner as in Example 1. It was measured.
As a result, the major axis direction was 500 μm or more and the minor axis direction 50 μm or more was 72%, and the major axis direction 1000 μm or more and the minor axis direction 100 μm or more was 68%.
（比較例２）
実施例１と同様の組成となるよう配合し、図８に示すようなＳＣ法の鋳造装置を用いて鋳造を行った。この水冷銅ロールの外径は４００ｍｍ、周速度は１ｍ／ｓとし、平均厚さ０．３０ｍｍのフレーク状の合金片を得た。
(Comparative Example 2)
The mixture was blended so as to have the same composition as in Example 1, and cast using an SC casting apparatus as shown in FIG. The water-cooled copper roll had an outer diameter of 400 mm, a peripheral speed of 1 m / s, and a flake-like alloy piece having an average thickness of 0.30 mm.
得られた合金片のＲ−リッチ相については、実施例１と同様の方法で短軸方向の長さが５μｍ以上、３μｍ以上であるＲ−リッチ相の面積率とＲ−リッチ相の平均間隔を測定した。（図１がＳＥＭの反射電子写真の一例である。図１で黒く見える部分は穴である。）
その結果、５μｍ以上の面積率は２％、３μｍ以上の面積率は５％であった。また、Ｒ−リッチ相の平均間隔は４．８μｍであった。
About the R-rich phase of the obtained alloy piece, the area ratio of the R-rich phase whose average length in the minor axis direction is 5 μm or more and 3 μm or more and the average interval between the R-rich phases in the same manner as in Example 1. Was measured. (FIG. 1 is an example of a reflection electrophotographic image of SEM. The portion that appears black in FIG. 1 is a hole.)
As a result, the area ratio of 5 μm or more was 2%, and the area ratio of 3 μm or more was 5%. The average interval between the R-rich phases was 4.8 μm.
ＳＣ合金の厚みの最大値は０．４８ｍｍであり、したがって長軸方向の長さ５００μｍ以上の柱状晶はなかった。なお、図２はこの合金片の断面の偏光顕微鏡写真の一例である。 The maximum value of the SC alloy thickness was 0.48 mm. Therefore, there were no columnar crystals with a length of 500 μm or more in the major axis direction. FIG. 2 is an example of a polarizing micrograph of the cross section of the alloy piece.
磁石の実施例
（実施例２）
実施例１で得られた合金塊を水素解砕、中粉砕、微粉砕の順に粉砕した。水素解砕工程の前工程である水素吸蔵工程の条件は、１００％水素雰囲気、大気圧で１時間保持とした。水素吸蔵反応開始時の金属片の温度は２５℃であった。また後工程である脱水素工程の条件は、１０Ｐａ真空中、５００℃で１時間保持とした。中粉砕にはブラウンミル装置を用い、水素解砕した粉末を１００％窒素雰囲気中で４２５μｍ以下まで粉砕した。この粉に、ステアリン酸亜鉛粉末を０．０７質量％添加し、１００％窒素雰囲気中においてＶ型ブレンダーで十分混合した後、ジェットミル装置で３．２μｍ（ＦＳＳＳ）まで微粉砕した。粉砕時の雰囲気は、４０００ｐｐｍの酸素を混合した窒素雰囲気中とした。その後、再度、１００％窒素雰囲気中においてＶ型ブレンダーで十分混合した。得られた粉体の酸素濃度は３１００ｐｐｍであった。またこの粉体の炭素濃度の分析から、粉体に混合されているステアリン酸亜鉛粉末は０．０５質量％であると計算された。
次に、得られた粉体を１００％窒素雰囲気中において横磁場中成型機でプレス成型した。成型圧は１１８ＭＰａであり、金型のキャビティ内の磁界は１２００ｋＡｍ−１とした。
得られた成型体を、１０-3Ｐａ真空中、５００℃で１時間保持し、次いで１０-3Ｐａ真空中、８００℃で２時間保持した後、１０-3Ｐａ真空中、１０６０℃で２時間保持して焼結させた。焼結密度は７．５×１０-3ｋｇｍ-3以上であり十分な大きさの密度となった。さらに、この焼結体をアルゴン雰囲気中、５４０℃で１時間熱処理した。
直流ＢＨカーブトレーサーでこの焼結体の磁気特性を測定した結果を表１に示す。
また、この焼結体の断面を鏡面研磨し、この面を偏光顕微鏡で観察したところ、結晶粒の大きさは平均で１０〜１５μｍであり、ほぼ均一の大きさであった。
Example of magnet (Example 2)
The alloy lump obtained in Example 1 was pulverized in the order of hydrogen pulverization, medium pulverization, and fine pulverization. The conditions of the hydrogen occlusion process, which is the previous process of the hydrogen crushing process, were maintained at 100% hydrogen atmosphere and atmospheric pressure for 1 hour. The temperature of the metal piece at the start of the hydrogen storage reaction was 25 ° C. The conditions for the dehydrogenation step, which is a subsequent step, were maintained at 500 ° C. for 1 hour in a 10 Pa vacuum. The medium pulverization was performed using a brown mill apparatus, and the hydrogen crushed powder was pulverized to 425 μm or less in a 100% nitrogen atmosphere. To this powder, 0.07% by mass of zinc stearate powder was added and sufficiently mixed with a V-type blender in a 100% nitrogen atmosphere, and then finely pulverized to 3.2 μm (FSSS) with a jet mill apparatus. The atmosphere during pulverization was a nitrogen atmosphere mixed with 4000 ppm of oxygen. Thereafter, the mixture was again sufficiently mixed with a V-type blender in a 100% nitrogen atmosphere. The oxygen concentration of the obtained powder was 3100 ppm. From the analysis of the carbon concentration of the powder, it was calculated that the zinc stearate powder mixed in the powder was 0.05% by mass.
Next, the obtained powder was press-molded in a transverse magnetic field molding machine in a 100% nitrogen atmosphere. The molding pressure was 118 MPa, and the magnetic field in the mold cavity was 1200 kAm −1 .
The resulting molded body, in 10 -3 Pa vacuum, and held 1 hour at 500 ° C., and then in a 10 -3 Pa vacuum, it was maintained for 2 hours at 800 ° C., in a 10 -3 Pa vacuum, 2 at 1060 ° C. It was held for a time and sintered. The sintered density was 7.5 × 10 −3 kgm −3 or more, which was a sufficiently large density. Furthermore, this sintered body was heat-treated at 540 ° C. for 1 hour in an argon atmosphere.
Table 1 shows the results of measuring the magnetic properties of this sintered body with a direct current BH curve tracer.
Moreover, when the cross section of this sintered compact was mirror-polished and this surface was observed with the polarization microscope, the average grain size was 10-15 micrometers, and it was a substantially uniform magnitude | size.
（比較例３、４）
比較例１で得られた合金塊および比較例２で得られた合金片をそれぞれ実施例２と同様の方法で粉砕して、３．２μｍ（ＦＳＳＳ）の大きさの粉体を得た。粉体の酸素濃度は３１００ｐｐｍであった。この粉体を使って、実施例２と同様の方法で磁場中成型、焼結し、異方性磁石を作製した。
(Comparative Examples 3 and 4)
The alloy ingot obtained in Comparative Example 1 and the alloy piece obtained in Comparative Example 2 were each pulverized in the same manner as in Example 2 to obtain a powder having a size of 3.2 μm (FSSS). The oxygen concentration of the powder was 3100 ppm. Using this powder, an anisotropic magnet was produced by molding and sintering in a magnetic field in the same manner as in Example 2.
得られた焼結体の磁気特性を表１に示す。
実施例２の保磁力（ｉＨｃ）は比較例３より１８５ｋＡｍ-1高い。この理由は、実施例１の合金塊ではプール状のＲ−リッチ相が少ないのに対し、比較例１の合金塊ではプール状のＲ−リッチ相が多くこのためＲ−リッチ相の分散状況が悪いためであると思われる。一方、実施例２の残留磁束密度（Ｂｒ）は比較例２より０．０２７Ｔ高い。これは配向率では２％高いことになる。この理由は、実施例１の合金塊の柱状晶は大きいのに対し、比較例２の合金片の柱状晶が小さいためであると思われる。
Table 1 shows the magnetic properties of the obtained sintered body.
The coercive force (iHc) of Example 2 is 185 kAm −1 higher than that of Comparative Example 3. The reason for this is that the alloy lump of Example 1 has a small amount of pool-like R-rich phase, whereas the alloy lump of Comparative Example 1 has a large amount of pool-like R-rich phase. It seems to be because it is bad. On the other hand, the residual magnetic flux density (Br) of Example 2 is 0.027 T higher than that of Comparative Example 2. This is 2% higher in the orientation ratio. This is probably because the columnar crystals of the alloy lump of Example 1 are large while the columnar crystals of the alloy pieces of Comparative Example 2 are small.
（実施例３〜１４）
表２に示す合金組成になるように、金属ネオジム、金属プラセオジム、金属ジスプロジウム、金属テルビウム、フェロボロン、コバルト、アルミニウム、銅、フェロニオブ、鉄、を配合し、実施例１と同様に溶解し、同様の鋳造装置にて鋳造を行った。なお表２に示すように、鋳型内面には８０Ｎｉ-２０Ｃｒの溶射膜、アルミナペーパーまたはアルミナ溶射膜を設けた。さらに、実施例３および実施例５では、合金の配合量を４３％増やすことで合金塊の厚みを厚くした。
それぞれの実施例で得られた合金塊の鋳型面は平滑であった。
(Examples 3 to 14)
Metal neodymium, metal praseodymium, metal dysprodium, metal terbium, ferroboron, cobalt, aluminum, copper, ferroniobium, and iron are blended so as to have the alloy composition shown in Table 2, and dissolved in the same manner as in Example 1. Casting was performed using a casting apparatus. As shown in Table 2, an 80Ni-20Cr sprayed film, alumina paper or alumina sprayed film was provided on the inner surface of the mold. Furthermore, in Example 3 and Example 5, the alloy lump thickness was increased by increasing the compounding amount of the alloy by 43%.
The mold surface of the alloy ingot obtained in each example was smooth.
それぞれの実施例で得られた合金塊のＲ−リッチ相について、実施例１と同様の方法で短軸方向の長さが５μｍ以上、３μｍ以上であるＲ−リッチ相の面積率とＲ−リッチ相の平均間隔を測定した。その結果を表２に示す。なお、α―Ｆｅと思われる相は無かった。 For the R-rich phase of the alloy ingot obtained in each example, the area ratio of the R-rich phase and the R-rich in the short axis direction of 5 μm or more and 3 μm or more in the same manner as in Example 1 The average spacing of the phases was measured. The results are shown in Table 2. There was no phase that seemed to be α-Fe.
また柱状晶についても、実施例１と同様の方法で、柱状晶の長軸方向の長さ５００μｍ以上、または１０００μｍ以上あり、短軸方向の長さが５０μｍ、または１００μｍ以上ある部分の面積率を測定した。その結果を表２に示す。 Also, for columnar crystals, the area ratio of the portion where the length in the major axis direction is 500 μm or more, or 1000 μm or more, and the length in the minor axis direction is 50 μm or 100 μm or more in the same manner as in Example 1. It was measured. The results are shown in Table 2.
（比較例５）
表３に示す合金組成になるように、金属ネオジム、金属プラセオジム、金属テルビウム、フェロボロン、コバルト、アルミニウム、銅、鉄、を配合し、比較例２と同様に溶解し、同様の鋳造装置にて鋳造を行い、平均厚さ０．３０ｍｍのフレーク状の合金片を得た。
(Comparative Example 5)
Metal neodymium, metal praseodymium, metal terbium, ferroboron, cobalt, aluminum, copper, and iron were blended so as to have the alloy composition shown in Table 3, dissolved in the same manner as in Comparative Example 2, and cast in the same casting apparatus. To obtain a flaky alloy piece having an average thickness of 0.30 mm.
得られた合金片のＲ−リッチ相について、実施例１と同様の方法で短軸方向の長さが５μｍ以上、３μｍ以上であるＲ−リッチ相の面積率とＲ−リッチ相の平均間隔を測定した。その結果を表３に示す。なお、α―Ｆｅと思われる相は無かった。 For the R-rich phase of the obtained alloy piece, the area ratio of the R-rich phase and the average interval between the R-rich phases with the length in the minor axis direction being 5 μm or more and 3 μm or more in the same manner as in Example 1. It was measured. The results are shown in Table 3. There was no phase that seemed to be α-Fe.
一方、合金片の厚みの最大値は０．４９ｍｍであり、したがって長軸方向の長さ５００μｍ以上の柱状晶はなかった。 On the other hand, the maximum value of the thickness of the alloy piece was 0.49 mm. Therefore, there was no columnar crystal having a length of 500 μm or more in the major axis direction.
磁石の実施例
（実施例１５）
実施例１３で得られた合金塊を実施例２と同様の方法で粉砕して、３．２μｍ（ＦＳＳＳ）の大きさの粉体を得た。粉体の酸素濃度は３１００ｐｐｍであった。この粉体を使って、実施例２と同様の方法で磁場中成型、焼結し、異方性磁石を作製した。
Example of magnet (Example 15)
The alloy lump obtained in Example 13 was pulverized in the same manner as in Example 2 to obtain a powder having a size of 3.2 μm (FSSS). The oxygen concentration of the powder was 3100 ppm. Using this powder, an anisotropic magnet was produced by molding and sintering in a magnetic field in the same manner as in Example 2.
直流ＢＨカーブトレーサーでこの焼結体の磁気特性を測定した結果を表４に示す。
また、この焼結体の断面を鏡面研磨し、この面を偏光顕微鏡で観察したところ、結晶粒の大きさは平均で１０〜１５μｍであり、ほぼ均一の大きさであった。
Table 4 shows the results of measuring the magnetic properties of this sintered body using a direct current BH curve tracer.
Moreover, when the cross section of this sintered compact was mirror-polished and this surface was observed with the polarization microscope, the average grain size was 10-15 micrometers, and it was a substantially uniform magnitude | size.
（比較例６）
比較例５で得られた合金片を実施例２と同様の方法で粉砕して、３．２μｍ（ＦＳＳＳ）の大きさの粉体を得た。粉体の酸素濃度は３１００ｐｐｍであった。この粉体を使って、実施例２と同様の方法で磁場中成型、焼結し、異方性磁石を作製した。
(Comparative Example 6)
The alloy piece obtained in Comparative Example 5 was pulverized in the same manner as in Example 2 to obtain a powder having a size of 3.2 μm (FSSS). The oxygen concentration of the powder was 3100 ppm. Using this powder, an anisotropic magnet was produced by molding and sintering in a magnetic field in the same manner as in Example 2.
この焼結体の断面を鏡面研磨し、この面を偏光顕微鏡で観察したところ、結晶粒の大きさは平均で１０〜１５μｍであり、ほぼ均一の大きさであった。 When the cross section of this sintered body was mirror-polished and this surface was observed with a polarizing microscope, the average size of the crystal grains was 10 to 15 μm, which was almost uniform.
一方、直流ＢＨカーブトレーサーでこの焼結体の磁気特性を測定した結果を表４に示す。Ｔｂを５．０質量％使用している比較例６の磁石と、Ｄｙを７．２質量％使用している実施例１５の磁石ではほぼ同等の磁気特性となった。 On the other hand, Table 4 shows the results of measuring the magnetic properties of the sintered body with a direct current BH curve tracer. The magnet of Comparative Example 6 using 5.0% by mass of Tb and the magnet of Example 15 using 7.2% by mass of Dy exhibited substantially the same magnetic characteristics.
本来なら、全希土類成分を一定として、保磁力が変わらない程度までＴｂをＤｙで置換するとＢｒが低下する。ところが本発明の合金から製作した磁石では配向率が高くなるため、保磁力が変わらない程度までＴｂをＤｙで置換してもＢｒ低下を抑えることができる。
なお実施例１５では比較例６のＴｂの全量をＤｙで置換したが、要求特性上や磁石製作工程上の制約からＴｂの全量をＤｙで置換できなくても、一部は置換できる。以上より、本発明の合金を使用することにより、資源が希少で非常に高価なＴｂの全量または一部を、Ｔｂよりかなり安価なＤｙで置換することができるので、磁石の価格を下げることができる。
Originally, if all rare earth components are constant and Tb is replaced with Dy to such an extent that the coercive force does not change, Br decreases. However, since the orientation ratio is high in the magnet manufactured from the alloy of the present invention, the decrease in Br can be suppressed even if Tb is replaced with Dy to the extent that the coercive force does not change.
In Example 15, the total amount of Tb in Comparative Example 6 was replaced with Dy. However, even if the total amount of Tb cannot be replaced with Dy due to restrictions on required characteristics and magnet manufacturing processes, a part of the Tb can be replaced. From the above, by using the alloy of the present invention, the total amount or a part of Tb, which is scarce and very expensive, can be replaced with Dy which is considerably cheaper than Tb, so that the price of the magnet can be reduced. it can.
本発明のＲ−Ｔ−Ｂ系焼結磁石用合金塊はハードディスク、磁気共鳴映像法、各種モーター等の磁石部品として利用できる。 The alloy block for an RTB-based sintered magnet of the present invention can be used as a magnetic part such as a hard disk, magnetic resonance imaging, and various motors.
１ 溶解チャンバー
２ 鋳造チャンバー
３ るつぼ
３１ 溶湯
４ 円筒状回転鋳型
Ｌ 円筒状鋳型の回転軸
５ 回転体（回転受け容器）
Ｒ 回転体（回転受け容器）の回転軸
６ 湯道
７ 合金塊
８ 鋳型駆動機構
９ 回転体（回転受け容器）の回転駆動機構
１０ 回転体の回転モーター
１１ 回転体（回転受け容器）の側面の穴部
２１ るつぼ
２２ タンディッシュ
２３ 冷却ロール
２４ ストリップ状合金片
DESCRIPTION OF SYMBOLS 1 Melting chamber 2 Casting chamber 3 Crucible 31 Molten metal 4 Cylindrical rotating mold L Rotating shaft of cylindrical mold 5 Rotating body (rotating receptacle)
R Rotating shaft of rotating body (rotating container) 6 Runner 7
Claims (17)
An RTB-based sintered magnet manufactured using the alloy ingot according to any one of claims 1 to 12 as a raw material.
Priority Applications (1)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2005110368A JP4879503B2 (en) | 2004-04-07 | 2005-04-06 | Alloy block for RTB-based sintered magnet, manufacturing method thereof and magnet |
Applications Claiming Priority (3)
| Application Number | Priority Date | Filing Date | Title |
|---|---|---|---|
| JP2004112810 | 2004-04-07 | ||
| JP2004112810 | 2004-04-07 | ||
| JP2005110368A JP4879503B2 (en) | 2004-04-07 | 2005-04-06 | Alloy block for RTB-based sintered magnet, manufacturing method thereof and magnet |
Publications (2)
| Publication Number | Publication Date |
|---|---|
| JP2005320628A true JP2005320628A (en) | 2005-11-17 |
| JP4879503B2 JP4879503B2 (en) | 2012-02-22 |
Family
ID=35468094
Family Applications (1)
| Application Number | Title | Priority Date | Filing Date |
|---|---|---|---|
| JP2005110368A Active JP4879503B2 (en) | 2004-04-07 | 2005-04-06 | Alloy block for RTB-based sintered magnet, manufacturing method thereof and magnet |
Country Status (1)
| Country | Link |
|---|---|
| JP (1) | JP4879503B2 (en) |
Cited By (7)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| WO2011004894A1 (en) | 2009-07-10 | 2011-01-13 | インターメタリックス株式会社 | Ndfeb sintered magnet, and process for production thereof |
| JP2013089687A (en) * | 2011-10-14 | 2013-05-13 | Nitto Denko Corp | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| WO2013146781A1 (en) | 2012-03-30 | 2013-10-03 | インターメタリックス株式会社 | NdFeB-BASED SINTERED MAGNET |
| US8562756B2 (en) | 2008-01-11 | 2013-10-22 | Intermetallics Co., Ltd. | NdFeB sintered magnet and method for producing the same |
| WO2014034650A1 (en) * | 2012-08-27 | 2014-03-06 | インターメタリックス株式会社 | Ndfeb-based sintered magnet |
| WO2015068681A1 (en) * | 2013-11-05 | 2015-05-14 | 株式会社Ihi | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| EP3534381A4 (en) * | 2016-10-28 | 2020-07-08 | IHI Corporation | Rare-earth permanent magnet and method for manufacturing rare-earth permanent magnet |
Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2000343193A (en) * | 1999-06-02 | 2000-12-12 | Kubota Corp | Centrifugal casting method and casting speed indicating device |
| JP2002301554A (en) * | 2000-08-31 | 2002-10-15 | Showa Denko Kk | Centrifugal casting method, centrifugal casting apparatus and alloy produced with this apparatus |
| JP2003077717A (en) * | 2001-09-03 | 2003-03-14 | Showa Denko Kk | Rare-earth magnetic alloy agglomeration, its manufacturing method and sintered magnet |
| JP2003334643A (en) * | 2002-03-11 | 2003-11-25 | Showa Denko Kk | Method for manufacturing rare earth alloy, alloy lump for r-t-b magnet, r-t-b magnet, r-t-b bonded magnet, alloy lump for r-t-b exchangeable spring magnet, r-t-b exchangeable spring magnet and r-t-b exchangeable spring bonded magnet |
-
2005
- 2005-04-06 JP JP2005110368A patent/JP4879503B2/en active Active
Patent Citations (4)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| JP2000343193A (en) * | 1999-06-02 | 2000-12-12 | Kubota Corp | Centrifugal casting method and casting speed indicating device |
| JP2002301554A (en) * | 2000-08-31 | 2002-10-15 | Showa Denko Kk | Centrifugal casting method, centrifugal casting apparatus and alloy produced with this apparatus |
| JP2003077717A (en) * | 2001-09-03 | 2003-03-14 | Showa Denko Kk | Rare-earth magnetic alloy agglomeration, its manufacturing method and sintered magnet |
| JP2003334643A (en) * | 2002-03-11 | 2003-11-25 | Showa Denko Kk | Method for manufacturing rare earth alloy, alloy lump for r-t-b magnet, r-t-b magnet, r-t-b bonded magnet, alloy lump for r-t-b exchangeable spring magnet, r-t-b exchangeable spring magnet and r-t-b exchangeable spring bonded magnet |
Cited By (19)
| Publication number | Priority date | Publication date | Assignee | Title |
|---|---|---|---|---|
| US8562756B2 (en) | 2008-01-11 | 2013-10-22 | Intermetallics Co., Ltd. | NdFeB sintered magnet and method for producing the same |
| US10854380B2 (en) | 2008-01-11 | 2020-12-01 | Daido Steel Co., Ltd. | NdFeB sintered magnet and method for producing the same |
| JP2015122517A (en) * | 2009-07-10 | 2015-07-02 | インターメタリックス株式会社 | Neodymium-iron-boron sintered magnet, and method for manufacturing the same |
| US9589714B2 (en) | 2009-07-10 | 2017-03-07 | Intermetallics Co., Ltd. | Sintered NdFeB magnet and method for manufacturing the same |
| WO2011004894A1 (en) | 2009-07-10 | 2011-01-13 | インターメタリックス株式会社 | Ndfeb sintered magnet, and process for production thereof |
| JP2013089687A (en) * | 2011-10-14 | 2013-05-13 | Nitto Denko Corp | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| WO2013146781A1 (en) | 2012-03-30 | 2013-10-03 | インターメタリックス株式会社 | NdFeB-BASED SINTERED MAGNET |
| US10546673B2 (en) | 2012-08-27 | 2020-01-28 | Intermetallics Co., Ltd. | NdFeB system sintered magnet |
| WO2014034650A1 (en) * | 2012-08-27 | 2014-03-06 | インターメタリックス株式会社 | Ndfeb-based sintered magnet |
| JPWO2015068681A1 (en) * | 2013-11-05 | 2017-03-09 | 株式会社Ihi | Rare earth permanent magnet and method for producing rare earth permanent magnet |
| AU2014344917B2 (en) * | 2013-11-05 | 2018-03-22 | Ihi Corporation | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| KR101936174B1 (en) * | 2013-11-05 | 2019-01-08 | 가부시키가이샤 아이에이치아이 | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| CN109887697A (en) * | 2013-11-05 | 2019-06-14 | 株式会社Ihi | The manufacturing method of rare earth permanent-magnetic material and rare earth permanent-magnetic material |
| EP3067900A4 (en) * | 2013-11-05 | 2017-08-02 | IHI Corporation | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| US10629343B2 (en) | 2013-11-05 | 2020-04-21 | Ihi Corporation | Rare earth permanent magnet and rare earth permanent magnet manufacturing method |
| WO2015068681A1 (en) * | 2013-11-05 | 2015-05-14 | 株式会社Ihi | Rare earth permanent magnet and method for manufacturing rare earth permanent magnet |
| CN109887697B (en) * | 2013-11-05 | 2021-07-20 | 株式会社Ihi | Rare earth permanent magnet material and method for producing rare earth permanent magnet material |
| EP3534381A4 (en) * | 2016-10-28 | 2020-07-08 | IHI Corporation | Rare-earth permanent magnet and method for manufacturing rare-earth permanent magnet |
| US11264154B2 (en) | 2016-10-28 | 2022-03-01 | Ihi Corporation | Rare earth permanent magnet and rare earth permanent magnet manufacturing method |
Also Published As
| Publication number | Publication date |
|---|---|
| JP4879503B2 (en) | 2012-02-22 |
Similar Documents
| Publication | Publication Date | Title |
|---|---|---|
| JP4832856B2 (en) | Method for producing RTB-based alloy and RTB-based alloy flakes, fine powder for RTB-based rare earth permanent magnet, RTB-based rare earth permanent magnet | |
| JP5949776B2 (en) | R-T-B type alloy flake and method for manufacturing R-T-B sintered magnet | |
| US7431070B2 (en) | Rare earth magnet alloy ingot, manufacturing method for the same, R-T-B type magnet alloy ingot, R-T-B type magnet, R-T-B type bonded magnet, R-T-B type exchange spring magnet alloy ingot, R-T-B type exchange spring magnet, and R-T-B type exchange spring bonded magnet | |
| JPWO2009075351A1 (en) | R-T-B type alloy and method for producing R-T-B type alloy, fine powder for R-T-B type rare earth permanent magnet, R-T-B type rare earth permanent magnet | |
| JP3267133B2 (en) | Alloy for rare earth magnet, method for producing the same, and method for producing permanent magnet | |
| JP4879503B2 (en) | Alloy block for RTB-based sintered magnet, manufacturing method thereof and magnet | |
| JP2008231535A (en) | R-t-b based alloy, method for producing r-t-b based alloy, fine powder for r-t-b based rare earth metal permanent magnet, and r-t-b based rare earth metal permanent magnet | |
| JP4389427B2 (en) | Sintered magnet using alloy powder for rare earth-iron-boron magnet | |
| JPH1036949A (en) | Alloy for rare earth magnet and its production | |
| JP4479944B2 (en) | Alloy flake for rare earth magnet and method for producing the same | |
| JP3234741B2 (en) | Alloy for rare earth magnet and method for producing the same | |
| JP4689652B2 (en) | Method for evaluating metal structure of RTB-based magnet alloy | |
| JP4366360B2 (en) | Raw material alloy for RTB-based permanent magnet and RTB-based permanent magnet | |
| JP3561692B2 (en) | Structure control method for rare earth element-containing alloy, alloy powder and magnet using the same | |
| JP4818547B2 (en) | Centrifugal casting method, centrifugal casting apparatus and alloy produced thereby | |
| WO2003020993A1 (en) | Rare earth magnet alloy ingot, manufacturing method for the same, r-t-b type magnet alloy ingot, r-t-b type magnet, r-t-b type bonded magnet, r-t-b type exchange spring magnet alloy ingot, r-t-b type exchange spring magnet, and r-t-b type exchange spring bonded magnet | |
| JP4318204B2 (en) | Rare earth-containing alloy flake manufacturing method, rare earth magnet alloy flake, rare earth sintered magnet alloy powder, rare earth sintered magnet, bonded magnet alloy powder, and bonded magnet | |
| JP2003183787A (en) | Principal phase-based alloy for rare earth magnet, manufacturing method therefor, mixed powder for sintered rare earth magnet, and rare earth magnet | |
| JP4754739B2 (en) | Alloy ingot for rare earth magnet, method for producing the same, and sintered magnet | |
| JP3721831B2 (en) | Rare earth magnet alloy and method for producing the same | |
| JP4040571B2 (en) | Method for producing rare earth-containing alloy flakes | |
| WO2019220950A1 (en) | Cast alloy flakes for r-t-b rare earth sintered magnet | |
| WO2009125671A1 (en) | R-t-b-base alloy, process for producing r-t-b-base alloy, fines for r-t-b-base rare earth permanent magnet, r-t-b-base rare earth permanent magnet, and process for producing r-t-b-base rare earth permanent magnet | |
| JP2003334643A (en) | Method for manufacturing rare earth alloy, alloy lump for r-t-b magnet, r-t-b magnet, r-t-b bonded magnet, alloy lump for r-t-b exchangeable spring magnet, r-t-b exchangeable spring magnet and r-t-b exchangeable spring bonded magnet | |
| JP3536943B2 (en) | Alloy for rare earth magnet and method for producing the same |
Legal Events
| Date | Code | Title | Description |
|---|---|---|---|
| A621 | Written request for application examination |
Free format text: JAPANESE INTERMEDIATE CODE: A621 Effective date: 20080304 |
|
| A977 | Report on retrieval |
Free format text: JAPANESE INTERMEDIATE CODE: A971007 Effective date: 20100311 |
|
| A131 | Notification of reasons for refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A131 Effective date: 20100316 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20100512 |
|
| A02 | Decision of refusal |
Free format text: JAPANESE INTERMEDIATE CODE: A02 Effective date: 20110614 |
|
| A521 | Written amendment |
Free format text: JAPANESE INTERMEDIATE CODE: A523 Effective date: 20110906 |
|
| A911 | Transfer to examiner for re-examination before appeal (zenchi) |
Free format text: JAPANESE INTERMEDIATE CODE: A911 Effective date: 20110913 |
|
| TRDD | Decision of grant or rejection written | ||
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 Effective date: 20111122 |
|
| A01 | Written decision to grant a patent or to grant a registration (utility model) |
Free format text: JAPANESE INTERMEDIATE CODE: A01 |
|
| A61 | First payment of annual fees (during grant procedure) |
Free format text: JAPANESE INTERMEDIATE CODE: A61 Effective date: 20111130 |
|
| R150 | Certificate of patent or registration of utility model |
Ref document number: 4879503 Country of ref document: JP Free format text: JAPANESE INTERMEDIATE CODE: R150 Free format text: JAPANESE INTERMEDIATE CODE: R150 |
|
| FPAY | Renewal fee payment (event date is renewal date of database) |
Free format text: PAYMENT UNTIL: 20141209 Year of fee payment: 3 |
|
| S111 | Request for change of ownership or part of ownership |
Free format text: JAPANESE INTERMEDIATE CODE: R313113 |
|
| R350 | Written notification of registration of transfer |
Free format text: JAPANESE INTERMEDIATE CODE: R350 |