JP2003188006A - Rare earth magnetic alloy sheet, its manufacturing method, sintered rare earth magnetic alloy powder, sintered rare earth magnet, metal powder for bonded magnet, and bonded magnet - Google Patents

Abstract

<P>PROBLEM TO BE SOLVED: To provide a rare earth magnet excellent in magnetic properties by suppressing fine R rich phase regions in a cast R-T-B system alloy ingot to obtain an alloy ingot having a very homogenized system for the purpose of evenly distributing the R rich phase in a magnet. <P>SOLUTION: The surface of a casting roll 3 is roughened finely so that a liquid alloy 1 cannot enter the roughened fine surface of the roll by its viscosity. As a result, the casting die surface side of an alloy 4 is not cooled quickly, and the growth of the fine R rich phase is suppressed. <P>COPYRIGHT: (C)2003,JPO

Description

Detailed Description of the Invention

[0001]

TECHNICAL FIELD The present invention relates to an RTB-based alloy (where R is at least one of rare earth elements including Y).
Seed, T is a transition metal that essentially requires Fe, and B is boron. Alloy flakes for rare earth magnets,
The present invention relates to an alloy powder for a rare earth sintered magnet, a rare earth sintered magnet, an alloy powder for a bonded magnet, and a bonded magnet.

[0002]

2. Description of the Related Art Recently, Nd-F has been used as an alloy for rare earth magnets.
The production amount of the e-B alloy is rapidly increasing due to its high characteristics, and it is used for HD (hard disk), MRI (magnetic resonance imaging), various motors, and the like. Usually, Nd is partially replaced with another rare earth element such as Pr or Dy, or Fe is partially replaced with another transition metal such as Co or Ni. −
They are collectively called R-T-B type alloys including B type alloys.
Here, R is at least one of rare earth elements including Y. In addition, T is a transition metal in which Fe is essential,
Part of Fe can be replaced with Co or Ni,
Cu, Al, Ti, V, Cr, Mn, N as additional elements
1 for b, Ta, Mo, W, Ca, Sn, Zr, Hf, etc.
You may add in combination with 1 type or multiple types. B is boron and can be partially substituted with C or N.

The R-T-B type alloy has a crystal composed of an R ₂ T ₁₄ B phase, which is a ferromagnetic phase contributing to the magnetization action, as a main phase,
It is a non-magnetic alloy in which an R-rich phase having a low melting point and enriched with rare earth elements coexists. Since it is an active metal, it is generally melted or cast in a vacuum or an inert gas. Further, in order to produce a sintered magnet from a cast RTB-based alloy ingot by a powder metallurgy method, the alloy ingot is crushed to about 3 μm (FSSS: measurement by a Fisher subsieve sizer) to obtain an alloy powder. After that, press molding in a magnetic field, sintering at a high temperature of about 1000 to 1100 ° C. in a sintering furnace, then heat treating and machining if necessary, further plating to improve corrosion resistance, and sintering. It is usually a magnet.

In the sintered magnet made of the R-T-B type alloy, the R-rich phase plays an important role as follows. 1) It has a low melting point and becomes a liquid phase at the time of sintering, which contributes to increasing the density of the magnet and thus improving the magnetization. 2) Eliminates grain boundary irregularities, reduces nucleation sites in the reverse magnetic domain, and increases coercive force. 3) Magnetically insulate the main phase to increase coercive force. Therefore, if the dispersion state of the R-rich phase in the molded magnet is poor, local sintering failure and deterioration of magnetism may occur, so it is important that the R-rich phase is uniformly dispersed in the molded magnet. Becomes Here, the distribution of the R-rich phase is greatly influenced by the structure of the R-T-B based alloy ingot when cast.

Another problem that occurs in casting R-T-B type alloys is that α-Fe is contained in the cast alloy ingot.
Is to be generated. α-Fe deteriorates the crushing efficiency when crushing an alloy lump, and if it remains in the magnet even after sintering, it deteriorates the magnetic properties of the magnet. Therefore, in the conventional alloy ingot, α-Fe was erased by performing homogenization treatment at high temperature for a long time as needed.

In the cast R-T-B type alloy ingot, α
In order to solve the problem of -Fe formation, a strip casting method (abbreviated as SC method) has been developed and used in the actual process as a method of casting an alloy ingot at a higher cooling rate. In the SC method, the molten metal is poured onto a water-cooled copper roll and a thin piece of about 0.1 to 1 mm is cast to rapidly solidify the alloy, which can suppress the precipitation of α-Fe. . Further, since the crystal structure of the alloy lump becomes finer, it becomes possible to produce an alloy having a structure in which the R-rich phase is finely dispersed. Thus, S
In the alloy cast by the C method, since the R-rich phase inside is finely dispersed, the dispersibility of the R-rich phase in the magnet after crushing and sintering is also good, and the magnetic characteristics of the magnet are improved. Has been successful. (JP-A-5-222488, JP-A-5-522
-295490 publication)

The alloy ingot cast by the SC method is also excellent in the homogeneity of the structure. The homogeneity of the structure can be compared by the crystal grain size and the dispersed state of the R-rich phase. In the alloy flakes produced by the SC method, chill crystals may occur on the casting roll side of the alloy flakes (hereinafter referred to as the mold surface side), but as a whole a moderately fine and homogeneous structure brought about by rapid solidification. Can be obtained.

As described above, the R-T-cast by the SC method
In the B-based alloy, the R-rich phase is finely dispersed and the production of α-Fe is also suppressed. Therefore, when producing a sintered magnet, the homogeneity of the R-rich phase in the final magnet is increased. In addition, it is possible to prevent the crushing and the adverse effect on magnetism due to α-Fe. Thus, the RTB cast by the SC method
The system alloy ingot has an excellent structure for producing a sintered magnet. However, as the properties of magnets have improved, it has become increasingly necessary to improve the homogeneity of the structure of the raw alloy ingot.

Therefore, for example, Japanese Patent Laid-Open No. 10-31711
No. 0 discloses a technique for producing a sintered magnet having good magnet characteristics by setting the area ratio of chill crystals on the mold surface side of a cast RTB-based alloy to 5% or less. It is disclosed. Since the chill crystal part becomes a fine powder having a particle size of 1 μm or less in the pulverizing step, it is believed that it disturbs the particle size distribution of the alloy powder and deteriorates magnetism.

[0010]

The present inventors have studied the relationship between the structure of a cast R-T-B type alloy ingot and the behavior at the time of hydrogen crushing or fine crushing, and as a result, the result of sintering It has been found that it is more important to control the dispersed state of the R-rich phase than the crystal grain size of the alloy lump in order to uniformly control the grain size of the alloy powder for magnets. And the volume ratio of chill crystals in the alloy ingot is actually several percent or less, which is more than the adverse effect of chill crystals.
It should be noted that a region where the dispersion state of the R-rich phase generated on the mold surface side in the alloy lump is extremely fine (fine R-rich phase region) has a greater effect on controlling the particle size of the magnet powder. I found it. That is, depending on the composition of the alloy ingot and the manufacturing conditions, R
Even when the amount of chill crystals in the TB alloy mass is reduced, the volume ratio of the fine R-rich phase region may exceed 50%, and the fine R-rich phase region has a particle size distribution of the magnet alloy powder. It was confirmed that it was necessary to reduce the fine R-rich phase region to improve the magnet characteristics.

Therefore, the present invention suppresses the formation of a fine R-rich phase region in a cast R-T-B type alloy ingot and produces an alloy ingot having a structure excellent in homogeneity.
It is an object of the present invention to provide a rare earth magnet having excellent magnet characteristics by making the R-rich phase distribution in the magnet uniform.

[0012]

The inventors of the present invention changed the casting conditions in the SC method, especially the surface condition of the rotating roll for casting, to form a fine R-rich phase region in the R-T-B type alloy flakes. The volume ratios were compared. Then, it was found that there is a relationship between the surface roughness of the alloy flakes on the mold surface side and the volume ratio of the fine R-rich phase regions. The present invention was made by the present inventors based on the above findings.

That is, the present invention comprises (1) an R-T-B type alloy (provided that R is at least one rare earth element including Y, T is a transition metal containing Fe as an essential component, and B is boron). In the alloy flakes for a rare earth magnet, the thickness is 0.1 mm or more and 0.5 mm or less, and the surface roughness of at least one surface of the alloy flakes is 10-point average roughness (R
z) is 5 μm or more and 50 μm or less, an alloy flakes for rare earth magnets. (2) The alloy flakes for a rare earth magnet according to (1) above, wherein the surface roughness of at least one surface of the alloy flakes is 10 μm or more and 25 μm or less in terms of ten-point average roughness (Rz). (3) The alloy flakes for a rare earth magnet as described in (1) or (2) above, wherein the volume ratio of the fine R-rich phase region in the alloy is 20% or less. (4) In a method for producing an alloy flakes for a rare earth magnet made of an R-T-B type alloy by a strip casting method, the surface roughness of the casting surface of a casting rotary roll is set to a ten-point average roughness (Rz).
And 5 μm or more and 100 μm or less. (5) In a method for producing an alloy flakes for a rare earth magnet made of an R-T-B type alloy by a strip casting method, the surface roughness of a casting surface of a casting rotary roll is set to a ten-point average roughness (Rz).
5 μm or more and 100 μm or less. The method for producing an alloy flakes for a rare earth magnet according to the above (1) to (3). (6) The alloy for rare earth magnets according to (4) or (5) above, wherein the surface roughness of the casting surface of the casting rotary roll is 10 μm or more and 50 μm or less in terms of ten-point average roughness (Rz). Method of manufacturing flakes. (7) An alloy powder for a rare earth sintered magnet, which is produced by subjecting the alloy flakes for a rare earth magnet according to (1) to (3) above to a hydrogen crushing step and then pulverizing with a jet mill. (8) A rare earth sintered magnet produced by powder metallurgy from the alloy powder for rare earth magnets according to (7). (9) An alloy powder for a bonded magnet manufactured by the HDDR method using the alloy flakes for a rare earth magnet described in (1) to (3) above. (10) A bond magnet manufactured using the alloy powder for a bond magnet according to the above (9). Is.

[0014]

BEST MODE FOR CARRYING OUT THE INVENTION N cast by a conventional SC method
FIG. 1 shows a backscattered electron image of a cross section of a thin piece of a d-Fe-B alloy (Nd 31.5 mass%) observed by an SEM (scanning electron microscope). In FIG. 1, the left side is the mold surface side and the right side is the free surface side. The surface roughness of the surface of the alloy flakes on the side of the mold is 3.4 μm in ten-point average roughness (Rz). The white part in FIG. 1 is the Nd-rich phase (the R-rich phase is called the Nd-rich phase because R is Nd), which is the free surface side (opposite to the casting surface side) from the center part of the alloy flakes. The surface on the side) extends in a lamellar shape in the thickness direction or forms a small pool having a directional shape such that the lamella is divided. However, on the template surface side, a region in which the Nd-rich phase is extremely finer and more randomly present than other regions is generated, and the present inventors have developed a region in which a fine R-rich phase region (R When the main component of Nd is Nd, it is also referred to as a fine Nd-rich phase region. This fine R-rich phase region usually starts from the mold surface side and extends toward the center. On the other hand, a portion where the fine R-rich phase region does not exist from the central portion to the free surface side is referred to as a normal portion here.

During the hydrogen disintegration process of the RTB-based alloy flakes during the production of sintered magnets, hydrogen is absorbed from the R-rich phase and expands into brittle hydride. Therefore, in the hydrogen disintegration, fine cracks along the R-rich phase or starting from the R-rich phase are introduced into the alloy. In the subsequent fine pulverization step, the alloy is broken due to a large amount of fine cracks generated by hydrogen crushing, so that the finer the R-rich phase dispersion in the alloy, the finer the grain size after fine pulverization tends to be. Therefore, the fine R-rich phase region has a stronger tendency to be finely divided than the normal portion. For example, in the case of alloy powder produced from the normal portion, the average particle size is about 3 μm as measured by FSSS (Fisher Subsieve Sizer). On the other hand, the alloy powder manufactured from the fine R-rich phase region has a high proportion of fine particles of 1 μm or less, and therefore the particle size distribution after fine pulverization becomes broad.

It is known that the dispersion state of the R-rich phase in the R-T-B type alloy can be controlled by controlling the cooling rate after the molten metal is solidified during casting or by heat treatment. Alternatively, JP-A-10-369
No. 49 publication. However, the cooling rate after solidification or the behavior of the change of the R-rich phase inside the fine R-rich phase region due to the heat treatment is difficult to control unlike the normal part, and the dispersion of the R-rich phase is less likely to become coarse and remains fine. Is.

The volume ratio of the fine R-rich phase region can be measured by the following method. FIG. 3 is a backscattered electron image of the same field of view as FIG. 1, but a line is drawn at the boundary between the fine R-rich phase region and the normal portion. Since the boundary between both regions can be easily determined from the dispersed state of the R-rich phase, the area ratio of the fine R-rich phase region in the field of view can be calculated using the image analysis device. The area ratio in the cross section corresponds to the volume ratio in the alloy. In the measurement of the volume ratio in the fine R-rich phase region, even if the alloy flakes were cast at the same time, the fine R
The change in the amount of the rich phase region is large between the flakes and within the same flakes. Therefore, the volume ratio of the fine R-rich phase region of the entire alloy is calculated by expanding the observation field of view at a low magnification of about 50 to 100 times, measuring about 5 to 10 thin pieces, and taking the average thereof. You can do it.

The RTB-based alloy flakes (Nd3) of the present invention
The backscattered electron beam image of the cross section of (1.5 mass%) is shown in FIG. In FIG. 2, the left side is the mold surface side and the right side is the free surface side. A feature of the alloy flakes of the present invention is that, in the flakes manufactured by the strip casting method, generation of fine R-rich phase regions is suppressed by controlling the surface roughness on the mold surface side. As shown in FIG. 2, in the alloy flakes of the present invention, the fine R-rich phase region does not exist on the mold surface side, and the R-rich phase dispersion state is extremely uniform from the mold surface to the free surface.

The relationship between the surface roughness of the surface of the alloy flakes produced by the strip casting method and the fine R-rich phase region can be explained as follows. In order for the alloy flakes to have a smooth surface on the mold surface side, the surface of the casting rotary roll must be smooth and have good wettability with the molten alloy. In such a state, the heat transfer from the molten metal to the mold is extremely good (the heat transfer coefficient is large), and the mold surface side of the alloy is excessively rapidly cooled. It is considered that the fine R-rich phase region has a strong tendency to be generated when the heat transfer coefficient of the mold and the molten metal is large and the mold surface side of the alloy is excessively rapidly cooled.

On the other hand, if fine irregularities are formed on the surface of the casting rotary roll, the molten metal cannot completely enter the fine irregularities on the surface of the casting rotary roll due to the viscosity of the molten alloy.
The non-contact portion is generated, and the heat transfer coefficient is reduced. As a result, it is considered that the mold surface side of the alloy is not excessively rapidly cooled and the generation of the fine R-rich phase region can be suppressed. When the surface roughness of the surface of the casting rotary roll is increased, the unevenness is transferred to the mold surface side of the alloy flakes to some extent, so that the surface roughness of the mold surface side of the alloy flakes naturally increases. It is presumed that the reason why the formation of the R-rich phase is suppressed by the alloy flakes whose surface on the mold surface side has an appropriate surface roughness is that excessive heat transfer when the molten metal is solidified is suppressed as described above. To be done.

However, if the surface roughness of the surface of the rotating roll for casting becomes excessively large, the molten metal will be able to enter the irregularities of the surface, and the heat transfer coefficient will increase again, and at the same time, the surface of the produced alloy flakes on the mold surface side. The roughness becomes even larger. In this case, the volume ratio of the fine R-rich phase region also increases again.

Even in the conventional SC method, alloy flakes having a homogeneous structure as shown in FIG. 2 were contained to some extent.
Since flakes containing a large amount of fine R-rich phase regions as shown in (3) are also generated at the same time, as a result, there is a problem in the homogeneity of the structure of the entire alloy. Such a conventional S
It is considered that the variation in the alloy structure produced by the C method is caused by the difference in the contact state between the roll surface and the molten metal, such as the subtle surface state of the rotating roll for casting, the supply state of the molten metal, and the atmosphere. On the other hand, in the present invention, since unevenness of an appropriate size is formed on the surface of the casting rotary roll, excessive heat transfer when the molten metal is solidified is eliminated, and the generation of the fine R-rich phase region is reproducibly suppressed. be able to. as a result,
It has become possible to manufacture alloy flakes having a uniform structure as shown in FIG. 2 in a high yield.

Further details of the present invention will be described. (1) Strip casting method The present invention relates to an RTB-based alloy flakes for rare earth magnets cast by the strip casting method. Here, casting of the R-T-B type alloy by the strip casting method will be described. FIG. 4 shows a schematic diagram of an apparatus for casting by the strip casting method. Usually R-T-B
Due to its active nature, the system alloys are melted with the refractory crucible 1 in a vacuum or an inert gas atmosphere. The molten metal of the melted alloy is kept at 1350 to 1500 ° C. for a predetermined time, and then, if necessary, through a tundish 2 provided with a rectifying mechanism and a slag removing mechanism, the inside of which is a water-cooled rotating roll for casting 3 Is supplied to. The supply rate of the molten metal and the rotation speed of the rotating roll are appropriately controlled according to the desired alloy thickness. Generally, the rotation speed of the rotating roll is about 1 to 3 m / s in terms of peripheral speed. Copper or a copper alloy is suitable as the material of the casting rotary roll because it has good thermal conductivity and is easily available. Depending on the material of the rotating roll and the surface condition of the roll, metal easily attaches to the surface of the rotating roll for casting, so if a cleaning device is installed as necessary,
The quality of the cast R-T-B type alloy is stable. The alloy 4 solidified on the rotating roll separates from the roll on the opposite side of the tundish and is collected in the collecting container 5. By installing a heating and cooling mechanism in this collection container, R
The state of rich phase tissue can be controlled.

The thickness of the alloy flakes of the present invention is preferably 0.1 mm or more and 0.5 mm or less. If the thickness of the alloy flakes is less than 0.1 mm, the solidification rate will increase excessively, the crystal grain size will become too fine, and it will be close to the finely pulverized grain size in the magnetizing step, so the orientation ratio and magnetization of the magnet will decrease. There is a problem of inviting. Further, when the thickness of the alloy flakes is more than 0.5 mm, the dispersibility of the Nd-rich phase decreases due to the decrease in solidification rate.
This causes problems such as precipitation of α-Fe.

(2) Surface Roughness of Cast Surface of Rotating Roll for Casting In the present invention, the strip casting method R-T-B is used.
When casting a magnet-based magnet alloy, the surface roughness of the casting surface of the rotating roll for casting is 10 μm or more in terms of ten-point average roughness (Rz) and is 10 μm or more.
It is set to 0 μm or less. Here, the surface roughness means JIS B
Measured under the conditions shown in 0601 "Definition and display of surface roughness", and the ten-point average roughness (Rz) is also defined therein. Specifically, first, a curve (roughness curve) obtained by removing a surface waviness component longer than a predetermined wavelength with a phase compensation type high-pass filter etc. from the cut (cross-sectional curve) when cut on a plane perpendicular to the measurement surface. Ask. From the roughness curve, the reference length is extracted in the direction of the average line, and from the average line of this extracted portion, the average value of the absolute values of the altitudes (Yp) of the 5th peak from the highest peak and the lowest absolute value. The sum of the absolute value of the absolute value (Yv) of the valley bottom to the fifth from the valley bottom is called the ten-point average roughness (Rz). Regarding the measurement parameters such as the reference length, standard values are specified by the above JIS for the surface roughness. The surface roughness of the alloy flakes on the mold surface side may fluctuate greatly, and at least 5 flakes should be measured and the average value should be used. Surface roughness is 5 μm
In the following, the effect of unevenness on the surface of the casting rotating roll cannot be obtained, and the contact with the molten metal is good, so that the heat transfer coefficient is large. As a result, it becomes easy to generate a fine R-rich phase region in the alloy. On the other hand, when the surface roughness of the casting rotary roll is 5 μm or more, the molten alloy cannot be completely entered into the fine irregularities on the surface of the rotary roll due to the viscosity of the molten alloy, resulting in a non-contact portion and heat. The transmission coefficient decreases. As a result, generation of fine R-rich phase in the alloy can be suppressed. The surface roughness is more preferably 10 μm or more in terms of ten-point average roughness (Rz).

The surface roughness of the rotating roll for casting is 100 μm.
If it exceeds, the depth of the irregularities on the surface of the rotating roll increases and the interval between the irregularities generally increases, so that the molten metal can enter the surface of the rotating roll without any gap.
Therefore, the heat transfer coefficient is likely to become excessively large again, and the fine R-rich phase region is likely to be generated in the alloy. Therefore, the surface roughness of the casting rotary roll is 100 μm or less, preferably 50 μm or less.

Surface Roughness of R-T-B Alloy Flakes In the present invention, at least one surface of the R-T-B alloy flakes for rare earth magnets has a ten-point average roughness (R).
z) is 5 μm or more and 50 μm or less. The surface on which the unevenness having the above-mentioned roughness is formed is the surface of the mold surface on which solidification begins during casting by the strip casting method, and the surface reflects the unevenness of the surface of the rotating roll. As described above, when the surface roughness of this surface is 5 μm or less or 50 μm or more, the volume ratio generated in the fine R-rich phase region is large, and the dispersion state of the R-rich phase in the alloy is nonuniform. As a result, the particle size distribution of the alloy powder after fine pulverization is widened in the manufacturing process of the sintered magnet, and the characteristics of the magnet are deteriorated, which is not preferable. In the present invention, the surface roughness of one surface of the alloy flakes is 5 μm or more and 50 μm or less, more preferably 7 μm or more and 25 μm or less.

Volume Ratio of Fine R-Rich Phase Region in Alloy In the present invention, the volume ratio of fine R-rich phase region in the R-T-B type alloy is 20% or less. As a result, the particle size distribution of the alloy powder after fine pulverization becomes narrow and uniform in the step of producing a sintered magnet, so that it is possible to obtain a homogeneous sintered magnet with no variation in characteristics.

Method for Producing Alloy Powder for Rare Earth Sintered Magnet, Rare Earth Sintered Magnet: The alloy flakes for a rare earth magnet made of the R—T—B type alloy cast according to the present invention are crushed, molded and sintered. It is possible to manufacture a high-performance anisotropic sintered magnet.

The crushing of the alloy flakes is usually carried out in the order of hydrogen crushing and fine crushing to produce an alloy powder of about 3 μm (FSSS). Here, the hydrogen disintegration is divided into a hydrogen storage step as a previous step and a dehydrogenation step as a subsequent step. In the hydrogen storage step, hydrogen is stored mainly in the R-rich phase of the alloy flakes in a hydrogen gas atmosphere with a pressure of 266 hPa to 0.3 MPa · G, and the R-rich phase is generated by the R-hydride formed at this time. Utilizing the volume expansion, the alloy flakes themselves are finely cracked or innumerable fine cracks are generated.
This hydrogen storage is carried out in the range of room temperature to 600 ° C., but in order to increase the volume expansion of the R-rich phase and efficiently divide it, the pressure of the hydrogen gas atmosphere is increased and room temperature to 100 ° C. It is preferable to carry out in the range of. The preferred processing time is 1 hour or more. The R-hydride produced in this hydrogen storage step is unstable in the atmosphere and is easily oxidized.
It is preferable to carry out a dehydrogenation treatment in which the alloy flakes are held in a vacuum of 1.33 hPa or less at about ° C. By this treatment, it is possible to change to an R-hydride that is stable in the atmosphere. A preferable treatment time for the dehydrogenation treatment is 30 minutes or longer. The dehydrogenation treatment can be omitted if the atmosphere is controlled to prevent oxidation in each process from hydrogen absorption to sintering.

The R-T-B type alloy flakes produced by the strip casting method of the present invention are characterized in that the R-rich phase is uniformly dispersed. The preferable average value of the R-rich phase interval depends on the crushed particle size in the magnet manufacturing process, but is generally 3 μm to 8 μm. In hydrogen disintegration, R
Cracks are introduced along the rich phase or starting from the R rich phase. Therefore, by crushing with hydrogen and then pulverizing, it is possible to maximize the effect of the R-rich phase uniformly and finely dispersed in the alloy, and to efficiently produce alloy powder with a very narrow particle size distribution. It is possible to produce. If a sintered magnet is produced without performing this hydrogen disintegration step, the characteristics of the produced sintered magnet will be inferior. (M. Sagawa et al. Proce
edingof the 5th internet
onal conferenceon Advance
d materials, Beijing China
(1999))

Fine pulverization means that the RTB-based alloy flakes are 3 μm.
It is to grind to about m (FSSS). As a pulverizing device for fine pulverization, a jet mill device is most suitable because it has good productivity and a narrow particle size distribution can be obtained.
By using the alloy flakes having a small amount of the fine R-rich phase region of the present invention, an alloy powder having a narrow particle size distribution can be produced with high efficiency and stability. The atmosphere for fine 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 in these inert gases. As a result, the pulverization efficiency is improved, and the oxygen concentration of the alloy powder after pulverization can be set to 1,000 to 10,000 ppm, and the alloy powder can be appropriately stabilized. At the same time, it is possible to suppress abnormal growth of crystal grains when the magnet is sintered.

When the above alloy powder is molded in a magnetic field,
A lubricant such as zinc stearate is preferably added to the powder in order to reduce the friction between the alloy powder and the inner wall of the mold and also reduce the friction between the powders to improve the orientation. The preferable addition amount is 0.01 to 1% by mass. The lubricant may be added before or after fine pulverization, but it is preferable to sufficiently mix it with 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.

The alloy powder pulverized to about 3 μm (FSSS) is pressed by a magnetic field molding machine. The mold is
It is manufactured by combining a magnetic material and a non-magnetic material in consideration of the magnetic field direction in the cavity. Molding pressure is 0.5-2t /
cm ² is preferred. The magnetic field in the cavity during molding is 5
20 kOe is preferred. The atmosphere at the time of molding is preferably an inert gas atmosphere such as argon gas or nitrogen gas, but in the case of the above-mentioned powder subjected to the oxidation resistance treatment, it is also possible in the air. Also, the molding is cold isostatic pressing (CIP: Col
d Isostatic Press) or a rubber mold for quasi-hydrostatic press (RIP: Rubber I)
Sostatic Press) is also possible. CI
Since P and RIP are hydrostatically compressed, the disorder of the orientation during molding is small, the orientation ratio can be increased more than that by die molding, and the maximum magnetic energy product can be increased.

Sintering of the molded body is performed at 1000 to 1100 ° C. The sintering atmosphere is preferably an argon gas atmosphere or a vacuum atmosphere of 1.33 × 10 ^-2 hPa or less. The holding time at the sintering temperature is preferably 1 hour or more. Further, upon sintering, it is necessary to remove as much hydrogen as possible from the lubricant and alloy powder in the molded body before reaching the sintering temperature. The preferable conditions for removing the lubricant are
A under vacuum of 1.33 × 10 ^-2 hPa or reduced pressure A
It is to hold at 300 to 500 ° C. for 30 minutes or more in an r-flow atmosphere. Further, the preferable conditions for removing hydrogen are:
700 to 90 in a vacuum of 1.33 × 10 ^-2 hPa or less
It is to hold at 0 ° C for 30 minutes or more.

After the completion of sintering, heat treatment can be performed at 500 to 650 ° C., if necessary, in order to improve the coercive force of the sintered magnet. A preferable atmosphere in this case is an argon gas atmosphere or a vacuum atmosphere, and a preferable holding time is 30 minutes or more.

The RTB-based alloy flakes for rare earth magnets produced in the present invention in which the generation of fine R-rich regions is suppressed can be suitably used for producing bonded magnets as well as sintered magnets. You can The case of producing a bonded magnet from the alloy flakes for rare earth magnets of the present invention will be described below.

The RTB-based alloy flakes of the present invention are first heat-treated if necessary. The purpose of heat treatment is α in the alloy
-Removal of Fe and coarsening of crystal grains. HDDR (Hydrogen) is used for producing alloy powder for producing bonded magnets.
ation Disproportionatio
n Desorption Recombinatio
n) is processed, but α-Fe existing in the alloy is HDD
It cannot be erased in the R treatment step, which causes a decrease in magnetism. Therefore, it is necessary to erase α-Fe before performing the HDDR process.

The average particle size of the alloy powder for bonded magnets is 50 to 300 μm, which is very large as compared with the alloy powder for sintered magnets. In the HDDR method, the crystal orientation of the original alloy coincides with the orientation of the recombined submicron crystal grains with a certain distribution. Therefore, two or more crystal grains having different crystallographic orientations in the alloy flakes of the raw material are contained in one bond magnet alloy powder, which means that the alloy powders include regions having greatly different crystallographic orientations, The orientation ratio of the magnet decreases, and the maximum magnetic energy product decreases. In order to avoid this, it is convenient that the crystal grain size in the alloy flakes is large. In an alloy cast by a rapid solidification method such as a strip casting method, the crystal grain size tends to be relatively small, so coarsening of the crystal grain by heat treatment is effective for improving the magnet characteristics.

Many reports have been made on the method for producing alloy powder for bonded magnets by the HDDR method (see, for example, T.W.
Takeshita et al, Proc. 10th
Int. Workshop on RE Magn
ets and theraplication,
Kyoto, Vol. 1 p551 (198
9)). The alloy powder is produced by the HDDR method as follows.

When the raw R-T-B type alloy flakes are heated in a hydrogen atmosphere, the R ₂ T ₁₄ B phase of the magnetic phase is mixed with α-Fe, RH ₂ and Fe ₂ B at about 700 to 850 ° C. Decompose into phases. Then, at a similar temperature, the atmosphere is switched to an inert gas atmosphere or a vacuum atmosphere to remove hydrogen, and the decomposed phase has an R ₂ T ₁₄ grain size of about submicron.
Rejoin phase B. At this time, if the composition of the alloy and the treatment conditions are appropriately controlled, the easy axis of magnetization of each recombined R ₂ T ₁₄ B phase (R ₂ T ₁₄ B phase C axis) becomes R _{2 in} the raw material alloy before decomposition. T
₁₄ Anisotropy magnet powder can be obtained which is substantially parallel to the C axis of the B phase and in which the directions of easy axes of magnetization of each fine crystal grain are aligned.

The alloy subjected to HDDR treatment is 50 to 30
The alloy powder may be pulverized to about 0 μm and then mixed with a resin to be compression-molded, injection-molded, etc., to obtain a bonded magnet.

The fine R-rich phase region has a strong tendency to be finely pulverized during the HDDR treatment as in the above-mentioned hydrogen disintegration treatment. The characteristics of the magnetic powder by the HDDR method decrease as the particle size decreases. Therefore, the RTB-based alloy of the present invention that suppresses the generation of the fine R-rich phase can be suitably used for producing magnetic powder for bonded magnets in HDDR treatment.

[0044]

EXAMPLES (Example 1) The alloy composition is Nd: 31.5% by mass, B: 1.00% by mass, Co: 1.0% by mass, A
l: 0.30% by mass, Cu: 0.10% by mass, and a raw material in which metal neodymium, ferroboron, cobalt, aluminum, copper, and iron are mixed so that the balance becomes iron, using an alumina crucible and argon gas. In an atmosphere of 1 atmosphere,
It was melted in a high frequency melting furnace, and the molten metal was cast by a strip casting method to prepare an alloy thin piece. The diameter of the rotating roller for casting is 300 mm, the material is pure copper, the inside is water-cooled, and the surface roughness of the casting surface is 20 μ in ten-point average roughness (Rz).
Adjusted to m. The peripheral speed of the roll during casting is 0.9 m / s
Then, alloy flakes having an average thickness of 0.30 mm were produced.

The surface roughness of the surface of the obtained alloy flakes on the mold surface was 10 μm in terms of ten-point average roughness (Rz). After embedding 10 alloy flakes and polishing, a backscattered electron image (BE) was taken for each alloy flakes with a scanning electron microscope (SEM).
I) was photographed at 100 times magnification. When the photographed photograph was taken into an image analyzer and measured, the volume ratio of the fine R-rich phase region was 3% or less.

(Example 2) The alloy composition was Nd 28.5.
%, B: 1.00% by mass, Co: 1.0% by mass, Al:
Using raw materials mixed so that 0.30% by mass, Cu: 0.10% by mass, and the balance being iron, casting was performed by the SC method under the same conditions as in Example 1 to produce an alloy flakes.

The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface was 10-point average roughness (R
z) is 9 μm, and the volume ratio of the fine R-rich phase region is
It was 3% or less.

(Comparative Example 1) Raw materials were blended in the same composition as in Example 1, and melting and casting by the SC method were carried out in the same manner as in Example 1. However, the surface roughness of the surface of the casting rotary roll was a ten-point average roughness (Rz) of 3.0 μm. The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface was 3.3 μ in ten-point average roughness (Rz).
m, and the volume ratio of the fine R-rich phase region was 41%.

(Comparative Example 2) Raw materials were mixed in the same composition as in Example 1, and melting and casting by the SC method were carried out in the same manner as in Example 1. However, the surface roughness of the surface of the rotating roll for casting was 120 μm in terms of ten-point average roughness (Rz). The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface was 86 μm in terms of ten-point average roughness (Rz).
The volume ratio of the fine R-rich phase region was 29%.

Next, an example of producing a sintered magnet will be described. (Example 3) The alloy flakes obtained in Example 1 were hydrolyzed and finely pulverized with a jet mill. The conditions for the hydrogen storage process, which is a process prior to the hydrogen disintegration process, are 100% hydrogen atmosphere, 2
It was kept at atmospheric pressure for 1 hour. The temperature of the metal piece at the start of the hydrogen storage reaction was 25 ° C. The condition of the dehydrogenation process, which is a post process, is 1 at 500 ° C. in a vacuum of 0.133 hPa.
It was supposed to hold time. To this powder, 0.07 mass% of zinc stearate powder was added, thoroughly mixed in a 100% nitrogen atmosphere with a V-type blender, and then finely pulverized with a jet mill device. The atmosphere at the time of pulverization was a nitrogen atmosphere mixed with 4000 ppm of oxygen. Then, it was thoroughly mixed again with a V-type blender in a 100% nitrogen atmosphere. The oxygen concentration of the obtained powder was 2500 ppm, and 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. As a result of measurement with a laser diffraction particle size distribution analyzer, the average particle size D50 is 5.10 μm, D10 is 2.10 μm, and D90 is 8.6.
It was 2 μm.

Next, the obtained powder was press-molded by a molding machine in a transverse magnetic field in a 100% nitrogen atmosphere. Molding pressure is 1.2
t / cm ² and the magnetic field in the mold cavity is 15 kO
e. The obtained molded body is 1.33 × 10 ⁻⁵ hPa
In vacuum at 500 ° C. for 1 hour, then 1.33 ×
After holding at 800 ° C. for 2 hours in a vacuum of 10 ⁻⁵ hPa,
Further, it was held in a vacuum of 1.33 × 10 ⁻⁵ hPa at 1050 ° C. for 2 hours for sintering. Sintered density is 7.5 g / cm
The density was ³ or more, which was a sufficiently large density. Further, this sintered body was heat-treated in an argon atmosphere at 560 ° C. for 1 hour to produce a sintered magnet.

The results of measuring the magnetic properties of this sintered magnet with a DC BH curve tracer are shown in Table 1. Table 1 also shows the oxygen concentration and particle size of the fine powder of the raw material for the sintered magnet.

(Comparative Examples 3 and 4) The alloy flakes obtained in Comparative Examples 1 and 2 were pulverized in the same manner as in Example 3 to obtain fine powder. Further, a sintered magnet was produced through the same molding and sintering steps as in Example 3. However, since the fine powder obtained from the alloy flakes of Comparative Examples 1 and 2 became difficult to sinter,
The sintering temperature was increased by 20 ° C. The results of the sintered magnets using the alloy flakes of Comparative Examples 1 and 2 are Comparative Examples 3 and 4.

Table 1 shows the results of measuring the magnetic properties of these sintered magnets with a DC BH curve tracer. Table 1 also shows the oxygen concentration and particle size of the fine powder of the raw material of each sintered magnet.
Shown in.

[0055]

[Table 1]

As shown in Table 1, in Comparative Examples 3 and 4, D10 is smaller than that in Example 3, so that it is understood that the ratio of extremely fine powder smaller than about 1 μm is large.
Such very fine particles are easily oxidized, and Comparative Example 3,
In Example 4, the oxygen concentration of the fine powder is slightly higher than that in Example 3. The magnetic properties of the magnets of Comparative Examples 3 and 4 are lower than those of Example 3 mainly because sintering becomes difficult due to an increase in oxygen concentration and crystal grains become coarse due to an increase in sintering temperature of 20 ° C. it is conceivable that.

Next, an example of producing a bonded magnet will be described. (Example 4) Alloy composition is Nd 28.5%, B: 1.0
Raw materials were blended so that 0 mass%, Co: 10.0 mass%, Ga: 0.5 mass% and the balance iron, and alloy flakes were cast by the SC method under the same conditions as in Example 1. The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface was 9 μm in ten-point average roughness (Rz), and the volume ratio of the fine R-rich phase region was 3% or less. , Α-Fe was not included.

The above alloy flakes were placed in hydrogen at 1 atm for 820
After holding at 0 ° C. for 1 hour, an HDDR treatment of holding at the same temperature in vacuum for 1 hour was performed. The resulting alloy powder is 150 μm
It was crushed by a brown mill below, 2.5 mass% of epoxy resin was added, and a magnetic field of 1.5 T was applied to perform compression molding to obtain a bonded magnet. Table 1 shows the magnetic characteristics of the obtained bonded magnet.

(Comparative Example 5) Raw materials were mixed in the same composition as in Example 4, and melting and casting by the SC method were carried out in the same manner as in Comparative Example 1. The obtained alloy flakes were evaluated in the same manner as in Example 1. As a result, the surface roughness of the mold surface was 3.1 μm in ten-point average roughness (Rz), and the volume ratio of the fine R-rich phase region was 40%. Met.

Then, a bonded magnet was produced in the same manner as in Example 4. Table 1 shows the magnetic characteristics of the obtained bonded magnet.

It can be seen from Table 1 that the bonded magnets of Example 4 and Comparative Example 5 have excellent magnetic properties in Example 4. In Comparative Example 5, the volume ratio of the fine R-rich region is high,
It can be estimated that the magnetism is low because the amount of relatively fine particles of 50 μm or less after HDDR treatment or pulverization is large.

[0062]

The alloy flakes of the present invention have a small volume ratio in the fine R-rich region, and the homogeneity of the dispersed state of the R-rich phase in the alloy is better than that of the conventional SC material. Therefore, the sintered magnet manufactured from the alloy flakes and the bonded magnet manufactured by the HDDR method exhibit superior magnetic properties to the conventional one.

[Brief description of drawings]

FIG. 1 is a view showing a cross-sectional structure of an alloy flakes for a rare earth magnet containing a fine R-rich phase produced by a conventional SC method.

FIG. 2 is a view showing a cross-sectional structure of an alloy thin piece for a rare earth magnet according to the present invention.

3 is a diagram in which a line is drawn on a boundary between a fine R-rich region and a normal portion in the cross-sectional structure of FIG.

FIG. 4 is a schematic diagram of a casting device of a strip casting method.

[Explanation of symbols]

1 refractory crucible 2 tundish 3 Casting rotating rolls 4 alloy 5 Collection container

─────────────────────────────────────────────────── ─── Continuation of front page (51) Int.Cl. ⁷ Identification code FI theme code (reference) B22F 9/08 B22F 9/08 MH01F 1/053 H01F 1/04 HF term (reference) 4E004 DB02 DB14 LC01 TA03 4K017 AA04 BA06 BB01 BB05 BB06 BB12 CA03 CA07 DA04 EA03 EA08 EC02 4K018 BA18 BB01 BB04 BB06 BD01 KA45 5E040 AA04 BB03 BD01 CA01 NN06

Claims (10)

[Claims] 1. A rare earth magnet made of an RTB-based alloy (wherein R is at least one of rare earth elements including Y, T is a transition metal essentially containing Fe, and B is boron). A rare earth characterized in that the thickness of the alloy flakes is 0.1 mm or more and 0.5 mm or less, and the surface roughness of at least one surface of the alloy flakes is 5 μm or more and 50 μm or less in terms of ten-point average roughness (Rz). Alloy flakes for magnets. 2. The alloy flakes for rare earth magnets according to claim 1, wherein at least one surface of said alloy flakes has a ten-point average roughness (Rz) of 7 μm or more and 25 μm or less. 3. The volume ratio of the fine R-rich phase region in the alloy is 2
It is 0% or less, The alloy flakes for rare earth magnets of Claim 1 or 2 characterized by the above-mentioned. 4. A method for producing an alloy flakes for a rare earth magnet made of an R-T-B type alloy by a strip casting method, wherein the casting surface of a casting roll has a surface roughness of 10 μm in average roughness (Rz) of 5 μm. A method for producing an alloy flakes for a rare earth magnet, which is 100 μm or less. 5. A method for producing an alloy flakes for a rare earth magnet made of an RTB-based alloy by a strip casting method, wherein the surface roughness of the casting surface of a casting rotary roll is 10 μm in terms of ten-point average roughness (Rz). The method for producing the alloy flakes for rare earth magnets according to claim 1, wherein the thickness is 100 μm or less. 6. The alloy flakes for rare earth magnets according to claim 4, wherein the surface roughness of the casting surface of the casting rotary roll is 10 μm or more and 50 μm or less in terms of ten-point average roughness (Rz). Manufacturing method. 7. An alloy powder for a rare earth sintered magnet, which is produced by subjecting the alloy flakes for a rare earth magnet according to claim 1 to a hydrogen crushing step and then pulverizing with a jet mill. 8. A rare earth sintered magnet produced from the alloy powder for a rare earth magnet according to claim 7 by a powder metallurgy method. 9. An alloy powder for a bonded magnet, which is produced by the HDDR method, using the alloy flakes for a rare earth magnet according to claim 1. 10. A bond magnet produced by using the alloy powder for a bond magnet according to claim 9.