JPH10324958A - Hard magnetic alloy compacted body, its production and thin type hard magnetic alloy compacted body - Google Patents

Abstract

PROBLEM TO BE SOLVED: To produce a hard magnetic alloy compacted body excellent in material strength, easily tinnable and furthermore excellent in hard magnetic performance, to provide a method for producing it and to produce a thin type hard magnetic alloy compacted body. SOLUTION: This hard magnetic alloy compacted body is the one composed of an Fe series or FeCo series alloy contg., by atom, 4 to 20% element R of one or more kinds among rare earth elements and 2 to 20% B, in which the alloy with a structure in which fine crystalline phases having <=100 nm average grain size are precipitated by rapid cooling is crystallized or undergoes grain- growth under stress to form a mixed phase state of soft magnetic phases or semihard magnetic phases and hard magnetic phases, furthermore, anisotropy is imparted to the crystal axes of the hard magnetic phases, and its coercive force is regulated to >=1 kOe.

Description

DETAILED DESCRIPTION OF THE INVENTION

[0001]

BACKGROUND OF THE INVENTION 1. Field of the Invention The present invention relates to a compacted hard magnetic alloy which has excellent magnetic performance and can be used for motors, actuators, speakers, and the like, a method of manufacturing the same, and a compact compacted hard magnetic alloy.

[0002]

2. Description of the Related Art Conventionally, Sm-Co sintered magnets are generally used as magnet materials having performance superior to ferrite magnets.
Fe-Nd-B sintered magnets, Fe-Nd-B quenched magnets, and the like are known, and Fe-Sm-
Many studies have been made on new alloy magnets such as N-based magnets.

[0003]

However, in these magnet materials, Nd of 10 atomic% or more or Sm of 8 atomic% or more is required, and since a large amount of expensive rare earth elements is used, ferrite magnets are required. There is a drawback that the production cost is higher than that of the prior art. Ferrite magnets are less expensive than these rare earth magnets, but have insufficient magnetic properties. For this reason, there has been a demand for the appearance of a magnet material which is harder than ferrite magnets at low cost.

In view of the above circumstances, the present inventors have studied hard magnetic materials having excellent hard magnetic properties at low cost, and as a result, as described in Japanese Patent Application No. 8-68822, Fe, One or more elements of Co and Ni as main components, one or two or more of rare earth elements and one or two of Zr, Nb, Ta and Hf.
Containing at least 50% or more, preferably 60% or more of the structure is a fine crystalline phase having an average crystal grain size of 100 nm or less, and the remainder is an amorphous phase; a bcc-Fe as the microcrystalline phase, invented a hard magnetic material, characterized by mainly comprising Fe-B compound and / or Fe ₁₄ R ₂ B ₁ containing the solid solution elements.

[0005] However, the above-mentioned hard magnetic material is manufactured by spraying a molten metal onto a rotating drum and quenching it to form a ribbon, for example, or by squirting the molten metal into a cooling gas to quench the powder in a droplet state to form a powder. Since it is manufactured by a manufacturing method for forming a shape, it can be obtained only in the form of a thin strip or a powder, and as it is, for example, a bulk (bulk) magnet having a shape usable for a motor, an actuator, a speaker, and the like. Could not get.

In general, as a method of molding a powdery magnetic substance and processing it into a bulk, a method of mixing a magnetic substance powder with a resin binder such as rubber or plastic and molding the mixture by compression molding or injection molding has heretofore been employed. Performed, and the magnets produced by these methods are known as "bonded magnets"
Because of its high degree of freedom in shape, it is widely used for electronic components and the like. However, in these conventional bonded magnets, the intermediary of the binder between the hard magnetic materials intervenes, and the density of the magnet portion cannot be increased with respect to the entire volume. In addition, there is a problem that the material strength is weak because of containing the resin, and it is difficult to reduce the thickness.

Further, from the viewpoint of thinning, conventional Sm-
Looking at Co magnets, when this magnet is polished to a thickness of several mm or less and made thinner, it is extremely fragile and has a drawback that handling problems easily occur. Although a bonded magnet can be manufactured using an Sm-Co magnet, the density cannot be increased due to the inclusion of the resin as described above, and the excellent hard magnetic properties of the Sm-Co magnet are impaired. was there.

SUMMARY OF THE INVENTION The present invention has been made to solve the above-mentioned problems, and has as its object to provide a hard magnetic alloy compact having excellent material strength, easy thickness reduction, and excellent hard magnetic performance. And a method of manufacturing the same and a thin hard magnetic alloy compact.

[0009]

SUMMARY OF THE INVENTION The present invention has been made in view of the above circumstances, and contains 4 to 20 atomic% of an element R composed of one or more rare earth elements and 2 to 20 atomic% of B. An alloy composed of an Fe-based or FeCo-based alloy and having a microcrystalline phase having an average crystal grain size of 100 nm or less precipitated by quenching is crystallized or grown under stress, and a soft magnetic phase is formed in the microstructure. Alternatively, a mixed phase of a quasi-hard magnetic phase and a hard magnetic phase is formed and anisotropy is given to the hard magnetic phase, so that a coercive force of 1 k
Oe or more. Further, the content of the element R composed of one or more of the rare earth elements is 3 to 20 atomic%,
Is composed of an Fe-based or FeCo-based alloy containing 2 to 20 atomic%, and the amorphous phase is crystallized under stress in an alloy having a structure including an amorphous phase by quenching. A fine crystalline phase having a crystal grain size of 100 nm or less is precipitated, a mixed state of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase is formed, and the hard magnetic phase is given anisotropy. The coercive force may be 1 kOe or more.

In the hard magnetic alloy compact having the above-mentioned structure, the alloy crystallized or grown under stress may have a
The heat treatment may be performed at a temperature of 00 to 1000 ° C., and a fine crystalline phase having an average crystal grain size of 100 nm or less may be precipitated as a main phase in the structure. In the hard magnetic alloy compact described above, a bcc (body-centered cubic structure) -Fe phase or a bcc-FeCo phase and Fe-B containing a solid solution element
And a soft magnetic phase or a quasi-hard magnetic phase having a coercive force of at least 1 kOe in which an amorphous phase is precipitated,
Each of the hard magnetic phases having a coercive force of 1 kOe or more in which at least a simple substance of the Fe ₁₄ R ₂ B phase (where R represents one or more elements among the rare earth elements) is precipitated is 10 vol.
(Volume)% or more may be included. In the compacted hard magnetic alloy described above, an alloy having a structure in which an amorphous phase or a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by quenching is crystallized or grain-grown under stress and compacted. It may be characterized by being formed.

In the hard magnetic alloy compact described above, the alloy containing an amorphous phase and exhibiting hard magnetism when crystallized is formed by solidification utilizing a softening phenomenon that occurs during a crystallization reaction. May be characterized. In the hard magnetic alloy compact described above, the alloy may be heated under stress. In the hard magnetic alloy compact described above, an alloy represented by the following composition formula may be used. T _x M _y R _z B _w However, T is, Fe, Co, at least one element of Ni, R is one or more elements of the rare earth elements, M is
One or more elements of Zr, Nb, Ta, Hf, Ti, V, Mo, W, and B represents boron,
y, z, w are atomic%, 50 ≦ x, 0 ≦ y ≦ 15, 3 ≦ z ≦ 2
0, 2 ≦ w ≦ 20. Next, x, y, z, and w indicating the composition ratio in the composition formula are atomic%, and 80 ≦ x ≦ 93, 0.5 ≦
y ≦ 5, 3 ≦ z ≦ 10, and 3 ≦ w ≦ 7. Further, x, y, z, and w indicating the composition ratio in the above composition formula are atomic%, and 86 ≦
x ≦ 93, 0.5 ≦ y ≦ 3, 3 ≦ z ≦ 7, and 3 ≦ w ≦ 5.

Next, a hard magnetic alloy compact having the following composition formula may be used. _{T x M y R z B w} E v However, T is, Fe, Co, at least one element of Ni, R is one or more elements of the rare earth elements, M is
One or more elements of Zr, Nb, Ta, Hf, Ti, V, Mo, W, B represents boron, E represents Cr, Al,
Pt, Ru, Rh, Pd, Os, Ir, Cu, Ag, A
x, y, z, w, and v, which represent at least one element of u, Ga, and Ge, and indicate the composition ratio are atomic%, and 50 ≦ x,
0 ≦ y ≦ 15, 3 ≦ z ≦ 20, 2 ≦ w ≦ 20, 0 ≦ v ≦ 10
It is. Further, x, y, z, indicating the composition ratio in the composition formula,
w and v are atomic%, 80 ≦ x ≦ 93, 0.5 ≦ y ≦ 5, 3 ≦ z
≦ 10, 3 ≦ w ≦ 7, and 0 ≦ v ≦ 5. X, y, z, indicating the composition ratio in the composition formula,
w and v are atomic%, 86 ≦ x ≦ 93, 0.5 ≦ y ≦ 3, 3 ≦ z
≦ 7, 3 ≦ w ≦ 5, and 0.1 ≦ v ≦ 5. In the compacted hard magnetic alloy described above, Si is replaced with T element by 0.1.
It may be characterized by being added by 5 to 5 atomic%. In the hard magnetic alloy compact described above, at least Nd or Pr is contained in the rare earth element R contained in the alloy.
Is preferably contained. In the hard magnetic alloy compact described above, the relative density of the compact obtained by compacting the alloy is preferably 90% or more. The above-described hard magnetic alloy compact preferably has a residual magnetization of 90 emu / g or more. The hard magnetic alloy compact described above has a residual magnetization (Ir) with respect to a saturation magnetization (Is).
Is preferably 0.6 or more.

Next, the method of the present invention comprises the step of
3 to 20 atomic% of element R consisting of at least one species, and 2-2 of B
By rapidly cooling an Fe-based or FeCo-based alloy containing 0 atomic% to form a microcrystalline phase of the alloy having an average crystal grain size of 100 nm or less, the alloy is crystallized or grown under stress. Forming a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase in the structure and imparting anisotropy to the hard magnetic phase. It is. Furthermore, the method of the present invention is characterized in that the element R consisting of one or more of the rare earth elements is 3 to 20 atomic%, and B is 2 to 20 atomic%.
By rapidly cooling the Fe-based or FeCo-based alloy containing at% by atom to make the alloy into a structure composed of an amorphous phase, the alloy is crystallized from the amorphous phase under stress, and the average in the structure is obtained. At least a step of precipitating a fine crystalline phase having a crystal grain size of 100 nm or less and forming a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase, and imparting anisotropy to the hard magnetic phase It is characterized by comprising.

In the above-mentioned method for producing a hard magnetic alloy compact, the alloy is crystallized or grain-grown under stress and then subjected to a heat treatment at 400 to 1000 ° C. so that the structure has an average crystal grain size of 100 nm. The following fine crystalline phase may be precipitated as a main phase.
In the method for manufacturing a hard magnetic alloy compact described above, after quenching the alloy to form a structure in which an amorphous phase or a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated, the alloy is subjected to stress. The method may be characterized in that crystallization or grain growth is performed and consolidation is performed. In the method for producing a hard magnetic alloy compact described above,
A method characterized by solidifying and forming an alloy containing an amorphous phase and exhibiting hard magnetism when crystallized by utilizing a softening phenomenon that occurs during a crystallization reaction may be used. In the method for producing a hard magnetic alloy compact described above, the method may be characterized in that the alloy is heated under stress. Next, in the method of manufacturing a hard magnetic alloy compact described above, a method of obtaining a compact having a relative density of 90% or more by consolidating the alloy may be used.

Next, the thin hard magnetic alloy compact of the present invention is characterized in that an Fe-based or FeC alloy containing 3 to 20 atomic% of an element R composed of one or more rare earth elements and 2 to 20 atomic% of B is used.
o-based alloy, the ratio h / S of the thickness h to the area S is
The relationship of 0.01 ≦ h / S ^0.5 ≦ 0.3 may be satisfied, and the relative density may be 90% or more. Next, the element R composed of at least one of the rare earth elements is 3 to
20% by atom and Fe-based or FeCo-based alloy containing 2 to 20% by atom of B, having a thickness of 50 μm or more and 2 mm
And the relative density may be 90% or more. Furthermore, the average crystal grain size of 10
An alloy having a structure in which a fine crystalline phase of 0 nm or less is precipitated is crystallized or grain-grown under stress, and a mixed state of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase is formed in the structure. A thin hard magnetic alloy compact formed by being formed may be used. Further, the alloy having a structure including an amorphous phase by quenching is crystallized under stress to have an average crystal grain size of 100 nm in the structure.
The following fine crystalline phase may be precipitated, and a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase may be formed.

Next, a thin hard magnetic material having the characteristics of the compacted hard magnetic alloy described in any of the above, wherein the thickness is 50 μm or more and 2 mm or less and the relative density is 90% or more. A magnetic alloy compact may be used. A thin hard magnetic alloy compact obtained by the method for manufacturing a hard magnetic alloy compact according to any one of the above, having a thickness of 50 μm or more and 2 mm or less and a relative density of 90% or more. good.

[0017]

DESCRIPTION OF THE PREFERRED EMBODIMENTS The present invention will be described below in detail.
The compacted hard magnetic alloy according to the present invention basically contains 4 to 20 atomic% of the element R composed of at least one of the rare earth elements.
Alloy composed of a Fe-based or FeCo-based alloy containing 2 to 20 atomic% of B and having a microcrystalline phase having an average crystal grain size of 100 nm or less precipitated by rapid cooling, or an amorphous phase Are crystallized or grown under stress to form a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase in a structure in which a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated. Alternatively, the average crystal grain size of 100
A fine crystalline phase of nm or less is precipitated and the above-mentioned mixed phase state is formed. Preferably, the hard magnetic phase has magnetic anisotropy.
The Fe-based or FeCo-based alloy is an alloy containing an amorphous phase (amorphous alloy) or an amorphous alloy containing some crystalline phase, which exhibits hard magnetism when crystallized. Is used.

In order to manufacture such a hard magnetic alloy compact, first, an alloy powder (granules) for molding is prepared. This alloy powder is obtained by a step of rapidly cooling the amorphous alloy from a molten metal to obtain a thin strip or powder, and a step of pulverizing the thin strip into a powder. Among the particle diameters of the alloy powder obtained here, the particle diameter of the powder used in the subsequent step is preferably in the range of 35 μm to 150 μm, and more preferably in the range of 50 to 100 μm. The reason for this is that a powder having a large particle diameter exceeding 150 μm may not be sufficiently amorphized at the time of production, and a powder having a small particle diameter of less than 30 μm is pulverized by a mill or the like. This is because, when it is converted, foreign matter such as a problem of oxidation or a part of the constituent material of the inner wall of the mill or the crushing blade may be mixed. As a method for obtaining an amorphous alloy or an amorphous alloy containing some crystalline phase from the above-mentioned molten metal, a method of spraying the molten metal on a rotating drum and quenching it to form a thin ribbon, and blowing the molten metal into a cooling gas And quenching in the form of droplets to form a powder, or a method such as sputtering or CVD. The amorphous alloy used in the present invention may be prepared by any of these methods. It may be. The alloy ribbon or alloy powder obtained by quenching has a structure in which a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated, or a structure composed of an amorphous phase.

Then, the obtained alloy powder is compacted at a high pressure simultaneously with or subsequent to crystallization of the amorphous phase in the alloy powder or grain growth of the fine crystalline phase under stress. The above average crystal grain size 100 nm
A soft magnetic phase or a quasi-hard magnetic phase and a mixed phase of a hard magnetic phase are formed in a structure in which the following fine crystalline phase is precipitated, or an average crystal grain size in a structure composed of the amorphous phase. A fine crystalline phase of 100 nm or less is precipitated, and the mixed phase state is formed. More preferably, the hard magnetic phase is provided with anisotropy. When the anisotropy is given to the hard magnetic phase as described above, higher remanent magnetization (Ir) can be obtained as compared with the case of isotropic.

It is preferable to heat the alloy powder during crystallization or grain growth under stress. When the alloy powder is compacted, it is preferable to use a softening phenomenon that occurs during a crystallization reaction to perform solidification molding. Here, the solidification molding utilizing the softening phenomenon during the crystallization reaction of the amorphous alloy is performed when the amorphous phase in the amorphous alloy is heated at the crystallization temperature or at the previous stage. When such a softening phenomenon occurs, the powders of the amorphous alloys are pressed together under pressure and integrated, so that the softened amorphous alloy is solidified and formed to have a high density ( This is because a hard magnetic alloy compact having a high relative density) can be obtained. Further, when solidifying and molding by pressurization, it is preferable to use an alloy containing 50% by weight or more of an amorphous phase from the viewpoint that a strong bond can be obtained and a permanent magnet having strong hard magnetism can be obtained.

As a specific example of producing a compact using the above alloy powder, an amorphous alloy is formed by applying a pulse current to the alloy powder while applying pressure to the alloy powder using a discharge plasma sintering apparatus. By heating at or near the crystallization temperature for a predetermined time, crystallization or grain growth is performed to obtain a compact, or a pulse current is applied to the alloy powder and the temperature is increased to increase the temperature of the amorphous alloy. A compact may be obtained by applying pressure from above and below or right and left with a pressurizing body such as a punch at a temperature near the crystallization temperature.

FIG. 1 shows a main part of an example of a discharge plasma sintering apparatus suitable for producing a hard magnetic alloy compact according to the present invention. A die 1 made of cemented carbide such as WC, an upper punch 2 and a lower punch 3 made of cemented carbide such as WC inserted into the die 1, and a WC provided outside the die 1.
Supports the outer frame die 8 made of cemented carbide and the lower punch 3,
A base 4 serving as one electrode when a pulse current to be described later flows, a base 5 pressing the upper punch 2 downward and serving as the other electrode through which the pulse current flows, and being sandwiched between upper and lower punches 2 and 3. And a thermocouple 7 for measuring the temperature of the alloy powder 6.

FIG. 3 shows the overall structure of the plasma sintering apparatus. The plasma sintering apparatus A shown in FIG. 3 is a type of discharge plasma sintering machine called Model SPS-2050 manufactured by Sumitomo Coal Mining Co., Ltd., and has a structure shown in FIG. 1 as a main part. The apparatus shown in FIG. 3 has an upper base 11 and a lower base 12, and a chamber 13 is provided in contact with the upper base 11.
The chamber 13 is connected to a vacuum exhaust device and an atmosphere gas supply device (not shown), and is filled with the raw material powder filled between the upper and lower punches 2 and 3. 6 can be maintained in a desired atmosphere such as an inert gas atmosphere. 1 and 3, the energizing device is omitted, but an energizing device separately provided is connected to the upper and lower punches 2, 3 and the bases 4, 5, and as shown in FIG. A suitable pulse current can be supplied through the punches 2 and 3 and the bases 4 and 5.

In order to produce a target compact using the spark plasma sintering apparatus shown in FIGS. 1 and 3, for example, an alloy powder 6 is put between upper and lower punches 2 and 3, and a chamber 13 is formed. Is simultaneously evacuated and molded by applying pressure from above and below with the punches 2 and 3, and at the same time, for example, a pulse current as shown in FIG. By heating at a temperature in the vicinity for a predetermined time, crystallization or grain growth under stress produces a compact.

Here, the applied pressure for performing the discharge plasma sintering method is 200 to 1500 MPa, preferably 500 to 1500 MPa.
It is preferable that crystallization or grain growth is performed at 1000 MPa and the molding is performed. If the applied pressure is less than 200 MPa, it is difficult to impart anisotropy to the hard magnetic phase, and the porosity of the obtained compact is large, and the molding density is undesirably small. The applied pressure is 1500
If the pressure exceeds MPa, the strength of the WC die at a high temperature is insufficient, which is not preferable. However, if the die is made of an alloy having a higher strength and the pressurizing mechanism of the plasma sintering device is strengthened, it is higher. Of course, pressure may be used. Here, the heating rate when heating the alloy powder 6 is 1
0 ° C./min (0.17 ° C. min. Second) or more, preferably 20 ° C./min.
Min (0.33 ° C./min) or more. Heating rate is 10 ℃
If it is less than / min, the crystal grains become coarse, so that the magnetic exchange coupling force between the adjacent hard magnetic phases in the amorphous phase or the soft magnetic phase is weakened, and the hard magnetic properties are deteriorated.

[0026] When performing discharge plasma sintering method, and the sintering temperature and Ts, the crystallization starting temperature of the amorphous alloy when the _{Tx, T x -200 ℃ ≦ Ts} ℃ ≦ T x + 200 ℃ It is preferable to perform sintering in a temperature range that satisfies the following relationship. When the sintering temperature Ts is less than T _x -200 ° C., the temperature is too low, it becomes difficult to crystallize, no longer available softening phenomenon in the crystallization temperature near, undesirable because it can not produce a dense sintered body. Sintering temperature Ts exceeds T _x + 200 ° C. When fine crystalline phase degraded hard magnetic properties by grain growth, undesirable. In the discharge plasma sintering method using such a discharge plasma device, the temperature of the alloy powder can be quickly raised at a predetermined speed by an electric current, and
Since the temperature of the alloy powder can be strictly controlled according to the value of the energizing current, the temperature can be controlled much more accurately than, for example, heating by a heater, whereby sintering can be performed under predesigned conditions close to ideal conditions.

Further, in the above-mentioned example, the compact is produced by solidifying and molding the alloy powder simultaneously with or subsequent to the crystallization or grain growth of the alloy powder under stress by the spark plasma sintering method. The alloy powder is preferably filled into a mold and simultaneously or with crystallization or grain growth by heating to a temperature at or near the crystallization temperature of the amorphous alloy while pressing, for example, in a hot press. Subsequently, a compact may be produced by solidification molding.

When the alloy powder is solidified and formed during the onset of the softening phenomenon, the relative density of the compact is adjusted to 90% or more, more preferably 95% or more, by adjusting the pressure, temperature, and molding time. It is preferable to compact.
As a result, the obtained hard magnetic alloy compact becomes a strong sintered body having an extremely dense structure, and becomes a permanent magnet that is firm in physical properties and has a small size and strong hard magnetism.

After crystallization or grain growth of the alloy powder under stress, the alloy powder is simultaneously or successively compacted.
A fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated as a main phase in the compact by heat treatment at 00 to 1000 ° C. Thereby, hard magnetic properties are exhibited. If the heat treatment temperature (annealing temperature) is less than 400 ° C., the amount of the R ₂ Fe ₁₄ B phase that contributes to the hard magnetic properties is small, and sufficient hard magnetic properties cannot be obtained. On the other hand, if the heat treatment temperature exceeds 1000 ° C., the grain growth of the fine crystal phase occurs, and the hard magnetic characteristics are undesirably reduced. In particular, conditions were selected so that the fine crystal phase having an average crystal grain size of 100 nm or less was 60% by volume or more of the compact and the remainder was an amorphous phase. An Fe phase or a bcc-FeCo phase;
If a Fe ₁₄ R ₂ B phase (where R represents one or more of the rare earth elements) is formed, a compact having extremely high hard magnetic properties can be obtained.

Further, conditions are selected so that the soft magnetic phase having a coercive force of 1 kOe or less and the quasi-hard magnetic phase and the hard magnetic phase having a coercive force of 1 kOe or more are each contained in the compact at 10 vol (volume)% or more. Coercive force is 1 kOe
At least a bcc (body-centered cubic structure) -Fe phase or a bcc-FeCo phase, a Fe-B compound containing a solid solution element, and an amorphous phase are precipitated in the following soft magnetic phase or quasi-hard magnetic phase, In addition, if at least a simple substance of the Fe ₁₄ R ₂ B phase (where R represents one or more of the rare earth elements) precipitates in the hard magnetic phase having a coercive force of 1 kOe or more, the soft magnetic This is preferable in that it can have the features of the magnetic phase and the hard magnetic phase. When the soft magnetic phase having a coercive force of 1 kOe or less is less than 10 vol (volume)%,
Although the coercive force of the compact becomes large, it is not preferable because the concentration of the rare earth element necessary for constituting the hard magnetic phase becomes high. The hard magnetic phase having a coercive force of 1 kOe or more is 10 v
When the content is less than ol (volume)%, the coercive force of the compact of the hard magnetic alloy becomes small, which is not preferable.

The hard magnetic alloy compact obtained by the above method exhibits exchange coupling characteristics in which a fine soft magnetic phase and a hard magnetic phase obtained by realizing a fine structure are combined, and Since the Fe concentration is higher than that of the conventional rare earth magnet, the residual magnetization is 90 emu / g or more, and the ratio of the residual magnetization (Ir) to the saturation magnetization (Is) (square ratio Ir / Is) is 0.6 or more. A permanent magnet molded body can be obtained. Further, in the compact of the hard magnetic alloy, the amorphous magnetic powder is crystallized or grain-grown under stress, so that the hard magnetic phase has anisotropy. Get bigger,
It has a high remanent magnetization (Ir). In addition, this hard magnetic alloy compact is formed by bonding amorphous alloy powders under pressure to each other and integrating them, so that compared to conventional bonded magnets in which magnetic powders are bonded using a binder. It is a permanent magnet that is rigid in physical properties and small in size and has strong hard magnetism. Further, since the compacted hard magnetic alloy of the present invention is formed from powder as described above, it can be formed into various thin shapes.

The above-mentioned plasma sintering is performed, and the applied pressure is 200 to 1500 MPa, preferably 500 to 1000 MPa.
By crystallizing or growing the grains with MPa and solidifying and shaping, it can be easily processed into a thin plate having a sufficient density, for example, a thin plate having a relative density of 90% or more and a thickness in the range of 50 μm to 2 mm. it can. When an amorphous ribbon is obtained by the aforementioned quenching method, the thickness is about several tens μm to 100 μm.
The limit is about m, and the width also varies depending on the scale of the manufacturing equipment, but is usually about several mm to several tens mm. This is because a sufficient cooling rate is necessary to obtain an amorphous phase, and if a ribbon with a greater thickness or width is to be obtained, the liquid quenching device including the rotating drum becomes extremely large. In addition, the production cost increases significantly, or it becomes difficult to obtain an amorphous phase. Also, in the case of a magnet hardened with resin like a bonded magnet, the relative density of about 80% is the limit.

On the other hand, according to the manufacturing method using the plasma sintering method described above, a hard magnetic thin plate having a relative density of 90% or more and a thickness of about 50 μm to 2 mm can be easily obtained. Next, the relationship between the thickness and the area of the thin hard magnetic alloy compact is 0.01 ≦ h / S ^0.5 ≦ 0.3, where h is the thickness, S is the area, and the ratio is h / S ^0.5.
Is preferably satisfied. In the present invention,
Since plasma sintering enables compacting and thinning, such a compact hard magnetic alloy compact can be obtained. It also has a feature that it is difficult to crack when the consolidation is performed at a low temperature of about 400 to 700 ° C. The consolidation at 400 to 700 ° C. is a heating condition at a lower temperature than the condition of heating to about 800 to 1000 ° C. as a sintering and solidifying molding temperature of a general metal material. The amount of thermal expansion and contraction associated with heating and cooling is small, and the thermal shock associated with heating and cooling is small, so that cracks are unlikely to occur. Here, the rapidly quenched ribbon having a thickness of about several tens of μm of the above-described composition is easily crystallized by heat treatment and has a brittle property. Therefore 400-700
The fact that solidification can be performed at a relatively low temperature of
Since the thermal shock accompanying heating and cooling can be reduced as compared with a material obtained by solidifying and molding at a high temperature of up to 1000 ° C., the solidifying molding temperature can be relatively low,
This is one reason why a thin hard magnetic material can be obtained in the present invention. Even if plasma sintering is performed, if the thickness is less than 50 μm, the amount of material existing between the punches 2 and 3 decreases, and it becomes difficult to uniformly apply pressure to the material existing between the punches 2 and 3. It is difficult to increase the relative density.

From the above, the compacted hard magnetic alloy of the present invention is useful as a permanent magnet used in various devices such as a motor, an actuator, and a speaker, and can reduce the manufacturing cost. The thin hard magnetic alloy compact of the present invention is different from a conventional Sm-Co magnet or the like in that when it is thin, it tends to crack.
Since it is easy to reduce the thickness to about μm to 2 mm, it is possible to provide a thin, strong magnet that cannot be obtained with a conventional magnet material of this type.

Next, the amorphous alloy which can be used for producing the compacted hard magnetic alloy of the present invention will be described in detail. The hard magnetic material according to the present invention can be represented by the following composition formula. TxMyRzBw T in the above composition formula represents one or more elements of Fe, Co, and Ni. Since these elements T are the main components of the hard magnetic material according to the present invention and are elements responsible for magnetism, the composition ratio x of the element T is 50 atomic% or more. As the composition ratio x of the element T increases, the saturation magnetization (Is) increases accordingly. High remanent magnetization (Ir) of 90 emu / g or more
In order to realize, the saturation magnetization (Is) must be at least 1
30 emu / g is necessary, and the element T
Is preferably at least 80 atomic%. Further, in order to obtain good hard magnetic properties, the content is preferably 93 atomic% or less. In the hard magnetic material of the present invention, it is necessary that Fe is contained as at least a part of the element T.

M in the above composition formula is Zr, Nb, Ta,
One or more of Hf, Ti, V, Mo, and W are represented, and these elements M have high amorphous forming ability.
In the hard magnetic material according to the present invention, by adding the element M, an amorphous phase can be formed even when the element R (rare earth element) has a low concentration. When the composition ratio y of the element M is increased by the substitution of the element R, the residual magnetization (I
r) increases, but the coercive force (iHc) decreases and the hard magnetic properties change to soft magnetic properties. In addition, when the element M is increased by the substitution of the element T which is responsible for magnetism, the saturation magnetization (Is) and the residual magnetization (Ir) decrease. Therefore, in order to obtain good hard magnetic properties, the composition ratio y of the element M is 0 atomic% or more and 1 atomic% or more.
It is preferably in the range of 5 atomic% or less, and 0.5 atomic%.
More preferably, it is in the range of at least 5 atomic%. Further, it is more preferable that the content be 0.5 atomic% or more and 3 atomic% or less. Further, in order to easily form an amorphous phase, it is more preferable to add 1 atomic% or more.

R in the above composition formula is a rare earth element (Sc,
Y, La, Ce, Pr, Nd, Pm, Sm, Eu, G
d, Tb, Dy, Ho, Er, Tm, Yb, and L
u) represents one or more elements. An alloy containing an element R, Fe and B and having an amorphous main phase is 873 to 1173.
The intermetallic compound R ₂ Fe ₁₄ B that precipitates when heated at an appropriate temperature in the range of K (600 to 900 ° C.) provides the hard magnetic material of the present invention with excellent hard magnetic properties.
As the composition ratio z of the element R increases, the saturation magnetization (Ir) decreases accordingly. In order to obtain a high residual magnetization (Ir) of 90 emu / g or more, the saturation magnetization (Is) needs to be at least 130 emu / g, and in order to satisfy this, the composition ratio z of the element R is 20 atomic% or less. Desirably. The element R is an element that easily forms an amorphous phase. If the composition ratio z of the element R is too small, a favorable amorphous phase or fine crystal phase cannot be obtained. Atomic% or more is desirable. In order to achieve both high saturation magnetization (Ir) and coercive force (iHc), 1
It is good to be 0 atom% or less, more preferably 7 atom% or less. Further, when part or all of the element R is composed of Nd and / or Pr, higher hard magnetic properties can be obtained.

B in the above composition formula is an element that easily forms an amorphous phase. Further, a compound R ₂ Fe ₁₄ which precipitates when an amorphous phase containing the elements R, Fe and B is heat-treated at an appropriate temperature in the range of 873 to 1173 K (600 to 900 ° C.).
B imparts hard magnetic properties to the hard magnetic material of the present invention. In order to obtain a good amorphous phase or fine crystalline phase, the concentration of B is preferably at least 2 atomic%, more preferably at least 3 atomic%. As a result, the saturation magnetization (Is), the remanent magnetization (Ir), and the coercive force (iHc) are reduced. %, More preferably 5 atomic% or less. The amorphous phase containing Fe and B is 60%.
When heated to a suitable temperature in the range of 0 ° C to 900 ° C, F
The compound of eB is precipitated.

The hard magnetic material of the present invention includes Cr, A
1, Pt, Ru, Rh, Pd, Os, Ir, Cu, A
One or more elements E among g, Au, Ga, and Ge may be added, and the hard magnetic material in that case can be represented by the following composition formula. TxMyRzBwEv In this case, the composition ratio x of the element T responsible for magnetism is preferably at least 50 at%, more preferably at least 80 at% and at most 93 at% from the viewpoint of increasing the saturation magnetization (Is), and 90 emu / g or higher remanent magnetization (Ir)
And 86 to realize a high coercive force (iHc)
It is preferable that the content be in the range of not less than atomic% and not more than 93 atomic%.
The composition ratio y of the element M in the above composition formula is preferably from 0 atomic% to 15 atomic% in order to obtain good hard magnetic properties.
More preferably, it is in the range of 0.5 at% to 5 at%, and in order to realize high remanent magnetization (Ir) of 90 emu / g or more, it is preferable to be in the range of 1 at% to 3 at%. . In order to obtain a higher remanent magnetization (Ir), the composition ratio may be set to 0.5 atomic% or more and 1 atomic% or less.

The composition ratio z of the element R in the above composition formula is determined so as to impart excellent hard magnetic properties to the hard magnetic material of the present invention.
In order to obtain a good amorphous phase or fine crystalline phase, it is preferably 3 atomic% or more and 20 atomic% or less, more preferably 3 atomic% or less.
In the range of not less than 10 atomic% and not more than 90 emu /
g or more to realize a high remanent magnetization (Ir) of 3 g or more.
It is preferable that the content be in the range of not less than atomic% and not more than 7%. The composition ratio w of B in the above composition formula is desirably 3 atomic% or more in order to obtain a good amorphous phase or fine crystalline phase. In order to obtain good hard magnetic properties, the composition ratio w of B is preferably 20 atomic% or less, more preferably 7 atomic%.
Atomic% or less, more preferably 5 atomic% or less. When the amorphous phase containing Fe and B is heated to an appropriate temperature in the range of 600 ° C. to 900 ° C., an Fe—B compound is precipitated. By adding the element E, the corrosion resistance of the hard magnetic material can be improved, or the refinement of the crystal structure can be promoted. However, if the composition ratio v of the element E is too high, the hard magnetic properties deteriorate, so the composition ratio v of the element E is preferably 10 atomic% or less, more preferably 5 atomic% or less. In order to achieve a high remanent magnetization (Ir) of 100 emu / g or more, it is preferable not to add the element E.

In the hard magnetic material of the present invention, F
If Co is included in addition to e, the absolute value of the temperature coefficient of magnetization when used in a shape having a permeance coefficient of 2 or more, and the absolute value of the temperature coefficient of magnetization when used in a shape having a permeance coefficient of 10 or more. This is preferable in that the absolute value and the absolute value of the temperature coefficient of the coercive force can be reduced.
The reason is that if Co is contained in the element T, the Curie temperature rises, so that the temperature change of magnetization and coercive force becomes small, and the squareness ratio of magnetization becomes high, so that the temperature change of magnetic characteristics becomes small. And the Co becomes bcc-Fe
This is because the temperature change of the remanent magnetization is reduced because it is included in the phase. If the content of Co is too large, the magnetic properties are deteriorated. Therefore, the content of Co is preferably 50 atomic% or less, more preferably 0.5 atomic% or more and 30 atomic% or less, and still more preferably 0.5 atomic% or more and 20 atomic% or less. And it is preferable to set appropriately according to the composition of the alloy, heat treatment conditions and the like.

Further, in the hard magnetic material of the present invention, Si
Is added by the element T substitution, the magnetic properties, especially the coercive force (i
Hc), and the maximum magnetic energy product ((BH) max)
And the absolute value of the temperature coefficient of magnetization when used in a shape with a permeance coefficient of 2 or more, particularly the temperature coefficient of magnetization when used in a shape with a permeance coefficient of 10 or more Can be reduced in absolute value. If the addition amount of Si is too large, the composition ratio of the element T becomes low, so that the magnetic properties of the hard magnetic material are rather deteriorated. Therefore, the addition amount of Si is preferably from 0.5 atomic% to 5 atomic%, more preferably 0.5 atomic%. It is preferably in the range of not less than atomic% and not more than 3 atomic%, and is preferably set appropriately according to the composition of the alloy, heat treatment conditions, and the like. The hard magnetic material having improved coercive force (iHc) and temperature characteristics in this way is particularly suitably used as a magnet and a sensor for a small motor.

Examples of particularly preferred amorphous alloys for producing the compact of the present invention include, for example, Fe ₈₈ Pr.
_{7 B 5, Fe 86 Pr 7} Nb 2 B 5, Fe 86 Nd 7 Zr 2 B 5, F
e ₈₆ Nd ₉ B ₅ , Fe ₈₄ Pr ₁₁ B ₅ , Fe ₈₈ Pr ₅ Nb
_{2 B 5, Fe 88 Nd 5} Nb 2 B 5, Fe 86 Nd 7 Nb 2 B 5, F
e ₈₉ Pr ₄ Nb ₂ B ₅ , Fe ₈₉ Nb ₂ Nd ₄ B ₅ , Fe ₈₉ Nb
₂ Pr ₄ B ₅ , Fe ₉₀ Nb ₂ Nd ₅ B ₃ , Fe ₉₀ Nb ₂ Pr ₅ B
_{3, Fe 89 Nb 2 Nd 5} B 4, Fe 89 Nb 2 Pr 5 B 4 can be cited. When alloys having these compositions are used, a strong compact is formed by the heat and pressure treatment, and a bcc-Fe phase and a Fe ₁₄ R ₂ B phase are formed in the generated fine crystal phase, which is excellent in hard magnetic properties. A permanent magnet can be obtained.

[0044]

The present invention will be described more specifically with reference to the following examples. (Preparation of amorphous alloy) Amorphous alloys having various compositions shown in Table 1 below were prepared by the following method. First, ingots of alloys having the respective compositions are prepared by an arc melting method, and a melt of this alloy is sprayed onto a rotating Cu roll in an Ar atmosphere to obtain a quenched ribbon having a thickness of about 20 μm. Was. The obtained quenched ribbon was pulverized using a rotor speed mill to obtain an amorphous alloy powder having a particle size of 50 μm to 150 μm.

(Preparation of Compacted Body) The crystallization temperature T _x (° C.) of each of the obtained amorphous alloy powders was measured by DSC (differential scanning calorimetry). Next, when filling this powder into a mold and sintering by heating and pressing with a hot press,
The sintering pressure was 636 MPa, the sintering time was 8 minutes, and the sintering temperature T _S (° C.) was changed to form a compact.

(Measurement) With respect to the obtained compact sample, the relative density (%) and the magnetic properties of the compact, such as residual magnetization Ir (T), squareness ratio (Ir / Is), and coercive force iHc
(KOe) was measured. During the measurement, the saturation magnetization (I
What is called s) is an applied magnetic field of 1.5T to 5T (tesla)
Shows the maximum magnetization obtained when the magnetization curve is measured by adding. Here, the relative density (%) is the true density (about 7.
5 g / cm ³ ), and the residual magnetization Ir
(T) is given by the following equation: Ir (T) = 4π × 7.5 × relative density × Ir (emu /
g) / 10000. FIG. 4 shows the composition of each amorphous alloy, the sintering temperature T _S (° C.), the sintering pressure P _S (MPa), the heat treatment temperature (° C.), and the direction in which the sintering pressure (sintering pressure P _S ) is applied.
, The saturation magnetization Is (T), the residual magnetization Ir (T), the squareness ratio Ir / Is, and the coercive force iHc in the Z direction and the X direction and the Y direction orthogonal to the Z direction. (KO
Table 1 shows the measurement results of e), the maximum magnetic energy product (BH) max, and the density (g / cm ³ ).

[0047]

[Table 1]

From the results shown in Table 1, when the amorphous alloys are crystallized or grown by the method of the present invention and solidified and formed by using the indicated amorphous alloys, all are dense and excellent hard magnetic properties. It can be seen that a compact having the following formula was obtained.
Further, Ir and Ir / Is are relatively high in the Z direction, and (B
H) It can be seen that high values were obtained for all the samples. In other words, it can be seen that the crystallization or grain growth under stress makes the hard magnetic phase anisotropic and improves the hard magnetic properties.

FIG. 5 shows a DSC (differential scanning calorimetry) curve of the amorphous alloy ribbon and a TMA (Thermo Mechanical Anisotropy).
alysis) curve. FIG. 5 shows that Fe ₈₈ Nb ₂ Nd ₅
Samples of the amorphous alloy having a composition of B _5, is obtained by measuring the DSC curve (b) and TMA curve (a) at a heating rate of 0.33 ° C. / sec. In the DSC curve (b) of FIG. 5, an exothermic peak is observed at about 850 ° C. This indicates that a crystallization reaction of the bcc-Fe or FeB compound has occurred. Looking at the TMA curve (a), the temperature is about 20 ° C. lower than the temperature at which the crystallization reaction occurs.
From the temperature range of about 427 ° C., which is 0 ° C. lower, the elongation of the sample increases as the temperature rises. This indicates that the alloy is softened near the crystallization temperature. For comparison, a TMA curve of a crystalline alloy having a composition of Fe ₈₈ Nb ₂ Nd ₅ B ₅ and containing no amorphous phase is also shown in FIG. FIG. 5 (a) shows that the crystalline alloy containing no amorphous material does not show any softening of the sample.

When the amorphous alloy powder particles are pressurized in this softening temperature range, the softened powder particles are tightly pressed together and bonded to form a dense compact having few pores (cavities). I do. Therefore, a compact having a high relative density can be obtained. FIG. 6 shows that the sintering pressure was 636 MPa and the sintering time was 8
FIG. 2 is a schematic view of a micrograph of the structure of a compact sample when sintered at various temperatures and various changes. FIG. 6 (a)
Is a sintering temperature of 400 ° C., and FIG.
FIG. 6C is a schematic diagram of a structure photograph at a sintering temperature of 600 ° C. From the schematic diagram of this structure photograph, as the sintering temperature rises, a dense compact having few pores (black irregular shaped portions indicate pores in the schematic diagram of the structure photograph of FIG. 6) is obtained. It can be seen that a sufficiently dense compact is obtained when the temperature is 600 ° C. or higher.

When the amorphous alloy is heated to the crystallization temperature, at least a part of the amorphous phase is crystallized. In FIG.
The sintering pressure is fixed at 636 MPa, the sintering time is fixed at 8 minutes, the temperature is variously changed, and the patterns obtained by X-ray diffraction of the compact sample immediately after sintering are shown. In FIG. 7, pattern (a) has a sintering temperature of 400 ° C., pattern (b) has a sintering temperature of 500 ° C., pattern (c) has a sintering temperature of 600 ° C., and pattern (d) has a sintering temperature of 650. ℃,
Is shown. In these patterns, 2θ
= 44.5 °, the halo pattern X appears near bcc-
This indicates the presence of the Fe crystal phase.

From this figure, it can be seen that at the sintering temperature of 500 ° C. or lower shown in the patterns (a) and (b), almost no bcc-Fe crystals are formed, and the halo pattern is an amorphous phase. c) to 600 ° C to 650 ° C shown in (d).
At the sintering temperature, the bcc-Fe crystal phase and N
It can be seen that the d ₂ Fe ₁₄ B phase and the Fe—B compound phase were formed, and showed characteristics as a hard magnetic material.

FIG. 8 shows the density of the compact when the sintering pressure was 636 MPa, the sintering time was 8 minutes, and the sintering temperature was changed. From this figure, it can be seen that the density increases as the sintering temperature increases, and that a density of about 7.45 g / cm ³ or more is obtained at a temperature of 500 ° C. or more.

When this result is compared with the X-ray diffraction pattern shown in FIG. 7, at a temperature of 500 ° C. or less where the amorphous state is maintained, a sufficiently high relative density can be obtained in the solidified compact. Not. On the other hand, bcc-
In the temperature range where the Fe crystal phase is formed, a sufficiently high density is obtained. From this, when solidification molding under stress utilizing the softening phenomenon that occurs during the crystallization reaction and the previous stage, crystallization and pressure bonding occur at the same time by one heat and pressure, sintering sufficiently densely, It can be seen that a compact having both excellent physical properties and excellent hard magnetic properties can be obtained.

In FIG. 8, as a comparative example, a sample (amorphous powder) which was previously heated and crystallized at 750 ° C. in a powder state and then annealed for 3 minutes was prepared. To
As in the case of the example, the density measured for a sample solidified at a temperature of about 600 ° C. with a sintering pressure of 636 MPa and a sintering time of 8 minutes is shown. In the sample of the comparative example, which was once crystallized in a powder state and then solidified and formed under the same conditions as in the example, the density was about 6.70 g / cm ³ or less. It can be seen that a compact having a sufficiently dense structure cannot be obtained.

FIG. 9 shows sintering temperatures of 600 ° C. and 650 ° C.
Shows the density of the compact when the sintering pressure is changed in the range of 260 to 636 MPa. For comparison, a powder sample (crystalline powder) which was previously heated and crystallized at 750 ° C. in a powder state and then annealed for 3 minutes was subjected to a sintering temperature of 600 ° C. and a sintering pressure of 636 MPa.
FIG. 9 shows the compacting density of the compact that was solidified under the conditions of a and holding time of 3 minutes. From FIG. 9, in the case where a sample obtained by solidifying and molding an amorphous powder is used, the density of the compact (consolidated body) is increased by increasing the compacting pressure (sintering pressure), and the pressure of 500 MPa or more is obtained. It can be seen that the molded article solidified by the above has almost the true density (7.7 g / cm ³ is the true density). On the other hand, when a sample obtained by solidifying and molding a crystalline powder is used, 636M
Despite solidification under high pressure of Pa or more,
It can be seen that only a low molding density was obtained.

Further, as a preferred example of the method of manufacturing a compact according to the present invention, a sintered die having a punch diameter of 18 mm is used, and an Fe-Nd-Nb-B alloy is used. An example of manufacturing a compact and a measurement result of magnetic properties of the obtained compact will be described.

(Preparation of Quenched Amorphous Alloy Strip) Fe ₈₈ Nd
_An amorphous alloy having a composition of ₅ Nb ₂ B ₅ and Fe ₈₆ Nd ₇ Nb ₂ B ₅ was prepared by the following method. First, an ingot of an alloy having each composition was prepared by an arc melting method,
Using a quartz nozzle with a slit diameter of 0.3 × 14 mm, A
r about 20 μm by spraying a melt of this alloy onto a single roll of Cu rotating in a r atmosphere.
Quenched ribbons of thickness The melt quenching conditions were as follows. Charged ingot mass 15 to 20 g Ultimate vacuum 6 × 10 ³ Pa or less Ar atmosphere pressure 15 cmHg Blow pressure 0.4 kgf / cm ³ Roll rotation speed 4000 rpm Blow temperature 1450 ° C. The obtained quenched alloy forms a good ribbon shape. However, there is no problem because the pulverization is performed in the next step.

(Preparation of Powder) The obtained rapidly quenched alloy ribbon was pulverized by a rotor speed mill and classified. When the weight ratio of the powder having each particle size was examined, it was found that any of the ribbons was preferably ground into a powder having a particle size of about 37 to 105 μm. For Fe ₈₆ Nd ₇ Nb ₂ B ₅ , the particle size is 37 to 53 μm.
Of Fe ₈₈ Nd ₅ Nb ₂ B ₅ ,
105 μm powder was the most abundant. From this, Nd
It is considered that an alloy having a composition having a high concentration is more brittle than an alloy having a low Nd concentration and is easily pulverized uniformly.

As a comparative example, F of a composition containing no rare earth
An e- (Nb, Zr) -B-based amorphous alloy ribbon was similarly prepared and pulverized. The yield of a powder having a particle size of 53 to 105 μm was 10% or less. From this, Fe-Nd-
It can be seen that the Nb-B-based amorphous alloy is easier to pulverize than the Fe- (Nb, Zr) -B-based amorphous alloy.

(X-ray diffraction of powder) FIGS. 10 and 11
Are Fe ₈₈ Nd ₅ Nb ₂ B ₅ powders (FIG. 10) and Fe ₈₆ Nd ₇ Nb ₂ B ₅ powders (FIG.
It shows the result of X-ray diffraction of 1). A broad diffraction peak was observed at around 2θ = 50 ° in powders of any particle size, indicating that all powders formed an amorphous phase. As a comparative example, Fe ₈₄ containing no rare earth was used.
X-ray diffraction was performed on the powder obtained by similarly pulverizing the amorphous alloy ribbon having the composition of Nb ₇ B _9.
Diffraction lines of the crystalline phase were observed in powders of ５３53 μm. This may be caused by crystallization during pulverization or mixing from a pulverizer. From these facts, Fe-Nd-
It is recognized that the Nb-B-based amorphous alloy has an advantage that it can be easily and finely pulverized while maintaining the amorphous phase.

(Production of compact) Amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder and amorphous Fe ₈₆ Nd ₇ having a particle size of 37 to 105 μm
Nb ₂ B ₅ powder 1 was charged between the upper and lower punches 2 and 3 of the discharge plasma sintering apparatus shown in FIG. 3, as the atmosphere in the chamber 13 is 3 × 10 ^-3 Pa or less By evacuating and applying pressure from above and below with the punches 2 and 3 and simultaneously applying a pulse current to heat,
A compact was obtained by crystallization or grain growth and solidification. The sintering conditions were a pressure of 636 MPa, a heating rate of 1.8 K / sec, a sintering temperature of 773 to 873 K, and a holding time of 480 seconds (8 minutes). Thereafter, the obtained compact was heat-treated in an atmosphere of 1.3 × 10 ⁻³ Pa or less under the conditions of a heating rate of 3 K / sec, a heat treatment temperature of 823 K to 1073 K, and a holding time of 180 seconds (3 minutes). .

FIG. 12 shows an amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder (powder A) and an amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder (powder B). (For organization in which fine crystal grains of nm order are precipitated in the amorphous phase) Fe ₈₈ Nb ₂ Nd ₅ B ₅
The powder (powder C) was heated at a rate of 1.8 K / sec and a sintering temperature of 8
It is a figure which shows the time (second) at the time of sintering at 73K, the temperature of each sample, and the result of having measured distance X between dies (expansion amount). Here, the temperature of the sample is measured by a thermocouple 7 attached to the side surface of the die.
Is defined as the distance between the upper and lower punches 2 and 3 of the die as shown in FIG. As is clear from the results shown in FIG. 12, nano-crystallized Fe ₈₆ Nb ₂ Nd ₇ B ₅ powder (powder C)
Shows that X monotonously increases with the temperature rise, whereas amorphous Fe ₈₈ Nb ₂ Nd ₅ B ₅ powder (powder A) and amorphous Fe ₈₆ Nb ₂ Nd ₇ B ₅ powder (powder A) It can be seen that the sintering of the powder B) stops expanding at around 500 K for 240 seconds or conversely shrinks. It is composed of amorphous Fe ₈₈ Nb ₂ Nd ₅ B ₅ powder (powder A), amorphous Fe ₈₆ Nb
b ₂ Nd ₇ B ₅ powder (powder B) has a temperature of 22 ° C. near the crystallization temperature.
This is considered to be due to softening around 240 ° C. at 7 ° C. and an increase in the density of the molded body.

(Structure and Magnetic Properties of Compacted Body) FIG.
Are amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder and amorphous Fe ₈₆
Nd ₇ Nb ₂ B ₅ powder was pressurized at a pressure of 636 MPa and the temperature was raised at a rate of 1.
8 ° C./sec, sintering temperature 500-600 ° C., holding time 48
It is a figure which shows the compacting density of the compacted body obtained by solidifying in 0 second (8 minutes). For comparison, crystalline Fe ₈₈ N
The molding density of d ₅ Nb ₂ B ₅ powder solidified molded obtained compacted body at a sintering conditions described above also shown in Figure 13. FIG.
Amorphous As apparent from the results shown in Fe ₈₈ Nd ₅ N
The compacting density obtained by sintering b ₂ B ₅ powder and amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder at a relatively high temperature of 600 ° C. has a molding density of about 7.5 × 10 ⁻³ kg / m ^3. And a bulk material of almost true density was obtained, whereas crystalline Fe ₈₈ N
In the case where the d ₅ Nb ₂ B ₅ powder solidified molded at 600 ° C. is observed that the molding density is not obtained only low molding density and ^{6.6 × 10 -3 kg / m 3} . The reason why a high molding density is obtained only when the crystalline alloy powder is solidified and molded is considered to be because the amorphous alloy softens near the crystallization temperature.

In the solidification molding of the amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder, the molding density of the bulk material tends to decrease as the sintering temperature decreases.
In solidification molding of e ₈₈ Nd ₅ Nb ₂ B ₅ powder, the molding density is about 7.5 × 1 even when the sintering temperature is 500 ° C.
0 ^-3 kg / m ³ is obtained, and it is recognized that a high molding density can be obtained even at a low temperature. The amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ alloy and the amorphous F ₈₈ measured at a heating rate of 0.67 ° C./sec by DSC (differential scanning calorimetry).
The crystallization onset temperatures of the e ₈₆ Nd ₇ Nb ₂ B ₅ alloy were 619 ° C. and 643 ° C., respectively. As described above, it is considered that the reason why the density of the amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ alloy is increased at a relatively low temperature is that the crystallization temperature is low.

FIG. 14 shows the results obtained by the above-mentioned spark plasma sintering method.
Fe obtained₈₆Nd₇Nb_TwoB_FiveCompacted body and Fe₈₈N
d_FiveNb_TwoB_FiveX-ray diffraction results immediately after sintering of the compact
Each is shown. The identification of the sample phase here
X-ray diffractometer using Co-Kα radiation
Was. From the results shown in FIG.₈₈Nd_FiveNb_TwoB_Five
After the sintering, the compact was diffracted by the bcc-Fe phase.
(Shown by ○ in the figure), Fe_TwoDiffraction peak due to B phase
(Indicated by □ in the figure) and Fe₁₄Nd_TwoTime by phase B
Folding peak (indicated by ● in the figure), and Fe₈₆N
d₇Nb_TwoB_FiveIn the compact, the bcc-Fe phase
Folding peak, Fe_ThreeDiffraction peak due to phase B (indicated by △ in the figure)
You. ) And Fe₁₄Nd_TwoA diffraction peak due to the B phase is seen.
It can be seen that these mixed phases are formed.
In addition, with each amorphous alloy powder,
The diffraction intensity of the crystalline phase has increased, and crystallization has progressed.
You can see that it is. Also, Fe ₈₆Nd₇Nb_TwoB_Fivealloy
The powder is compacted at 500 ° C and 550 ° C
In addition, relatively broad diffraction lines were obtained,
It seems to be a mixed phase of a quality phase and an amorphous phase. Thus relatively
In the sample in which crystallization has not progressed, it is shown in FIG.
6.6-7.0 × 10^-3kg / m^ThreeLow molding density
Or not obtained. From the results of FIG. 13 and FIG.
High density when crystallizing alloy powder simultaneously with solidification molding
It is found that it is advantageous for the formation.

FIG. 15 shows that an amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder and an amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder were pressed at a pressure of 636 M.
Pa, heating rate 1.8 ° C / sec, sintering temperature 500-60
It shows the magnetic properties after heat treatment of a consolidated body obtained by solidifying at 0 ° C. for 480 seconds (8 minutes) for a holding time of 3 ° C./min, 750 ° C., and a holding time of 180 seconds. The magnetic properties of the sample here were measured using a VSM (vibrating sample magnetometer) and applying an applied magnetic field of 1.5 T to the sample (0.5 m thick).
m, width 1.5 mm, length 6 mm) in the longitudinal direction. From the results shown in FIG.
_{The 88} Nd ₅ Nb ₂ B ₅ compact has a small change in remanence and squareness even when the sintering temperature is increased, but the Fe ₈₆ Nd ₇ Nb ₂ B
_{(5) In the} consolidated body, the remanent magnetization, the squareness ratio and the coercive force tended to increase as the sintering temperature increased.
It can be seen that the magnetic properties are excellent at 00 ° C.

In order to clarify the difference in density,
Amorphous Fe ₈₈ Nd ₅ Nb ₂ B using a 0 mm sintered die
FIG. 16 shows the relationship between the magnetic properties and the sintering temperature when the _five powders were solidified at a sintering pressure of 636 MPa. FIG. 16 shows that when the sintering temperature is lowered, the saturation magnetization (Is), the residual magnetization (Ir), and the squareness ratio (Ir / Is) are all reduced. FIG. 16 shows the relationship between the magnetic properties and the sintering pressure when the amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder was solidified at a sintering temperature of 600 ° C. FIG. 16 shows that when the sintering pressure is reduced, the saturation magnetization (Is), the residual magnetization (Ir), and the squareness ratio (Ir / Is) are all reduced. The reduction in the molding density when the sintering pressure and the sintering temperature during sintering are reduced has already been shown in FIGS. FIG. 18 shows that the amorphous F
The relationship between the molding density and the magnetic properties when solidifying e ₈₈ Nd ₅ Nb ₂ B ₅ powder is shown. As shown in FIG. 18, the saturation magnetization (Is), the residual magnetization (Ir), and the squareness ratio (I
r / Is) are both reduced. From this, it can be said that to obtain high magnetic properties, it is important to increase the density of the compact.

FIG. 19 shows the Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact compacted at a pressure of 636 MPa, a rate of temperature rise of 1.8 ° C./sec, a sintering temperature of 600 ° C., and a holding time of 480 seconds (8 minutes). FIG. 4 shows the results of X-ray diffraction of a Fe ₈₆ Nd ₇ Nb ₂ B ₅ compact after heat treatment at 750 ° C. FIG. FIG.
From the results shown in FIG. 9, in the Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact,
After the heat treatment at 750 ° C., the diffraction peak due to the bcc-Fe phase (indicated by ○ in the figure), the diffraction peak due to the Nd ₂ Fe ₁₄ B phase (indicated by ● in the figure), and the diffraction peak due to the Fe ₃ B phase (in the figure) , Δ)), and Fe ₈₆ Nd ₇ N
In the b ₂ B ₅ compact, a diffraction peak due to the bcc-Fe phase and a diffraction peak due to the Nd ₂ Fe ₁₄ B phase are observed.
It can be seen that a mixed phase of a soft magnetic phase (bcc-Fe) and a hard magnetic phase (Nd ₂ Fe ₁₄ B) is formed in these compacts. X-ray diffraction pattern bcc (100)
The crystal grain size of the bcc-Fe phase determined from the half width of the diffraction line is the above Fe ₈₆ Nd ₇ Nb ₂ B ₅ compact and Fe ₈₈ Nd ₅ N
In the b ₂ B ₅ compact, approximately 20 nm and 30
nm, which was the same as the value for the ribbon alloy of the same composition. Further, using a high-resolution transmission electron microscope, the above 7
The structure of the Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact heat-treated at 50 ° C. was observed.
An Fe phase and a Nd ₂ Fe ₁₄ B phase having a particle size of 20 nm were observed.
From these results, it can be seen that even in the hard magnetic alloy compact, a nanocrystalline double phase structure (a structure in which a plurality of types of crystal grains in the order of nm are precipitated together with an amorphous phase) after heat treatment at 750 ° C. as in the case of the ribbon alloy. It can be seen that is formed.

FIG. 20 shows that the amorphous powder was compressed at a pressure of 636 MPa.
a, Fe ₈₈ Nd ₅ N solidified and formed at a temperature rising rate of 1.8 K / sec, a sintering temperature of 600 ° C., and a holding time of 480 seconds (8 minutes)
The magnetization curves of the b ₂ B ₅ compact and the Fe ₈₆ Nd ₇ Nb ₂ B ₅ compact after heat treatment at 750 ° C. are shown. FIG.
At 0, the solid line shows the magnetization curve of Fe ₈₈ Nd ₅ Nb ₂ B ₅ , and the broken line shows the magnetization curve of Fe ₈₆ Nd ₇ Nb ₂ B ₅ . In any of the compacts, as in the case of the magnetic material composed of a single phase, a magnetization curve without any step was obtained. From this, in the obtained compact, the fine soft magnetic phase and the hard magnetic phase are magnetically coupled to each other, exhibiting a magnetization curve like a hard magnetic material composed of only one hard magnetic phase, That is, it is understood that the exchange coupling magnet characteristics are obtained. Further, the Fe ₈₆ Nd ₇ Nb ₂ B ₅ compact and the Fe ₈₈ Nd ₅ N
For the b ₂ B ₅ compact, the residual magnetization Ir (T), squareness ratio (Ir / Is), coercive force iHc
(KOe), and the maximum magnetic energy product (BH) max
(KJ / m ³ ) was measured. The results are shown in Table 2 below. For comparison, Table 2 also shows the results of measuring the magnetic properties of the amorphous alloy ribbon having the same alloy composition as that of the above-mentioned compact after heat treatment. The heat treatment temperature was 750 ° C. at which the best hard magnetic characteristics were obtained. FIG. 21 shows amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder and crystalline Fe ₈₈ Nd ₅ Nb ₂ B ₅
The pressure of each powder is 636 MPa, and the heating rate is 1.8K.
FIG. 6 shows magnetization curves after heat treatment at 750 ° C., which is the optimum heat treatment temperature, for each sample (consolidated body) that was solidified and molded at a temperature of 750 ° C./sec, a sintering temperature of 600 ° C., and a holding time of 480 seconds (8 minutes). FIG. 21 shows that the amorphous Fe ₈₈ Nd ₅ Nb ₂
B ₅ powder solidified molded sample, crystalline Fe _{8 8} Nd ₅ Nb ₂
B ₅ powder solidified molded samples show magnetization curves not both seen a step, an amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅
The sample obtained by solidifying and molding the powder was crystalline Fe ₈₈ Nd ₅ Nb ₂ B
_{5 It} can be seen that the hard magnetic characteristics are superior to those of the sample obtained by solidifying and molding the powder.

[0071]

[Table 2]

From these results, the compacts having any composition have a coercive force (iHc) substantially equal to that of the ribbon. Further, the maximum magnetic energy product (BH) max of the compact is inferior to that of the ribbon, which is considered to be due to a decrease in the squareness ratio.

FIG. 22 to FIG. 23 show amorphous Fe ₈₈ Nd ₅ N
The b ₂ B ₅ powder and the amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder were subjected to a pressure of 636 MPa, a heating rate of 1.8 K / sec, and a sintering temperature of 50.
The temperature of the compacted body obtained by solidifying and molding at 0 to 600 ° C. and a holding time of 480 seconds (8 minutes) is 3 K / min, 62
It shows the magnetic properties after heat treatment at a temperature of 7 to 827 ° C. and a holding time of 180 seconds. From the results shown in FIGS. 22 and 23, the sintering temperature of the Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact was 600 ° C.
It can be seen that when the heat treatment temperature is 1023 K, the coercive force is high and the magnetic properties are excellent. In addition, Fe ₈₆ Nd ₇ Nb ₂ B
_{5 When the} sintering temperature is 600 ° C, the change in magnetic properties is small even if the heat treatment temperature is increased.
It can be seen that when the heat treatment temperature is 750 ° C. at 50 ° C. or 500 ° C., the coercive force is high and the magnetic properties are excellent.

In Table 2 above, the magnetic properties of the sintered compacts
Showed inferior results compared to the strip alloy,
Is self-demagnetized due to the thickness of the bulk material (sintered compact)
It seems that the world has not been corrected. Therefore, various compositions
Pressure of 636MPa, sintering temperature 600 ° C,
Obtained by solidification molding with a retention time of 480 seconds (8 minutes)
Each of the high-density compacts is 5 × 5 × 5 mm (or 4 ×
4 × 4mm) and heat-treated under optimal heat treatment conditions.
After that, measurement was performed using a pulse magnetizer in an applied magnetic field of 5T.
FIGS. 24 to 27 show demagnetization curves of magnetization when set.
(The demagnetizing field correction was performed.) FIG. 24 shows an amorphous
Quality Fe₉₀Nb_TwoNd_FiveB_FiveX of the sample obtained using the powder
FIG. 7 is a diagram showing demagnetization curves of magnetization in the directions of Y, Y, and Z.
You. FIG. 25 shows that amorphous Fe₈₉Nb_TwoNd_FourB _FiveUsing powder
Of the magnetization of the sample obtained in the X, Y and Z directions
It is a figure showing a curve. FIG. 26 shows that amorphous Fe₇₆Co_TenN
b_TwoNd₇B_FiveX direction, Y direction of sample obtained using powder
It is a figure which shows the demagnetization curve of magnetization of a direction and a Z direction. FIG.
Is amorphous Fe₈₄Nb_TwoNd₇B_FiveSi_TwoObtained using powder
The demagnetization curves of the magnetization of the sample in the X, Y, and Z directions
FIG. When solidifying the amorphous powder, Z
Pressure in the direction
Crystallization or grain growth.

Each of the compacts of each composition has a remanent magnetization of 0.8 T or more and a coercive force of about 2.5 kOe or more, exhibiting better hard magnetic properties than those measured by a VSM with an applied magnetic field of 1.5 T. It turned out to be showing. As is clear from FIGS. 24 to 27, it can be seen that a bulging curve is obtained in each of the compacts in the Z direction. The reason why the hard magnetic properties are improved in the Z direction is that uniaxial anisotropy can be added to the hard magnetic phase by precipitating the hard magnetic phase under pressure (stress).

The pressure (Ps) of the amorphous powder having a composition of Fe ₆₆ Co ₂₀ Nb ₂ Pr ₇ B ₅ was set at 636 MPa.
The maximum magnetic energy product ((BH) _max parallel) measured in the pressing direction (Z direction) at the time of molding of the obtained bulk sample
And the maximum magnetic energy product ((BH) _max vertial) measured in the direction (X or Y direction) perpendicular to the pressing direction during molding, and the maximum magnetic energy product ((BH) _max parallel)
Divided by the maximum magnetic energy product ((BH) _max vertial) "((BH) _max parallel) / ((BH) _max vertia
l) "is used as an index indicating the degree of anisotropy, and the dependency of the value on the sintering temperature (Ts) is shown in FIG. From the results shown in FIG. 28, it can be seen that the degree of anisotropy tends to increase as the sintering temperature (Ts) increases. This is because a hard magnetic phase (Fe ₁₄ R ₂₎ that crystallizes and grows under uniaxial stress during solidification molding.
It is considered that the degree of anisotropy increases as a result of the large volume ratio of (B ₁ phase). Anisotropy increases in a sample having a high pressure (Ps) during molding. This is due to the easy magnetization of the hard magnetic phase (Fe ₁₄ R ₂ B ₁ phase) that crystallizes and grows under uniaxial stress. It is presumed that the axis (c-axis) is oriented as the pressure during molding increases.

Next, the F method is performed under the same conditions as those described above at a plasma sintering temperature of 600 ° C. and a pressure of 636 MPa.
A hard magnetic alloy compact sample having a composition of e ₇₆ Co ₁₀ Nb ₂ Nd ₇ B ₅ was prepared. Other sample preparation conditions were the same as those of the sample having the same composition shown in Table 1, and the sample shape was a 2 mm square rectangular parallelepiped. When this sample was subjected to plasma sintering, a sample manufactured under the condition of a temperature rising rate of 3 ° C./min (0.05 ° C./sec) and a sample fabricated at a temperature rising rate of 110 ° C./min (1.83 ° C./sec) were used. The magnetic properties in each direction of the samples manufactured under the conditions were measured at room temperature and compared. In the following measurement directions, the Z direction is the direction in which the punch is pressed during plasma sintering, and the X and Y directions are directions orthogonal to the Z direction.

A sample having a heating rate of 110 ° C./min (0.05 ° C./sec) iHc Is Ir Ir / Is (BH) max (kOe) (T) (T) (MGOe) X direction: 3.65 1.49 0.94 0.6285 8.2025 Y direction: 3.64 1.49 0.94 0.6316 8.2815 Z direction: 3.66 1.50 0.96 0.6432 9.2551 Sample with heating rate of 110 ° C./min (0.05 ° C./sec) iHc Is Ir Ir / Is (BH) max (kOe) (T) (T) T) (MGOe) X direction: 4.11 1.48 0.96 0.648 8.8254 Y direction: 4.11 1.47 0.95 0.6487 8.7296 Z direction: 4.11 1.50 0.98 0.6532 9.7696

From the above results, it is clear that the rate of temperature increase during plasma sintering has a large effect on the hard magnetic properties, and a higher rate of temperature increase provides better hard magnetic properties in all directions. It is clear that. It was also found that a thin plate having sufficiently excellent hard magnetic properties can be obtained even when the sample shape is reduced to about 2 mm.

Next, a thin magnet was manufactured utilizing the fact that the alloy of the composition system of the present invention softens in the crystallization temperature range.
When preparing the samples shown in Table 1 having the composition of Fe ₈₈ Nb ₂ Nd ₅ B ₅ , Ts = 873 K (600 ° C.) and Ps = 6
Applying the conditions of 36 MPa, adjusting the interval between the punches of the plasma sintering apparatus, the diameter is 10 to 18 mm, and the thickness is about 800
A μm disk-shaped thin magnet was produced. In addition, Fe ₆₆ Co ₂₀
When preparing a sample having the composition of Nb ₂ Pr ₇ B ₅ , solidification and molding were carried out in the same manner, and 13 mm on a side and a thickness of 300 to 5 mm.
A sheet-shaped thin magnet of 00 μm was produced. Thickness 800
For a μm sample, the density is 7.48 g / cm ³ (relative density of about 97
%), And a density of 7.34 for a sample having a thickness of 300 to 500 μm.
７7.39 g / cm ³ (relative density: about 95%) could be obtained. From these manufacturing results, a high-density molded magnet thinner than 1 mm can be manufactured without performing mechanical processing such as polishing.

FIG. 29 shows the dependency of the composition of Fe ₆₆ Co ₂₀ Nb ₂ Pr ₇ B ₅ on the thickness of a bulk sample (compacted body) (area S = 162 mm ² ). Although the molding density is reduced by reducing the thickness, the thickness is 0.3 mm (300 μm).
(H / S ^0.5 = 0.016), a high relative density of 95 to 96% was obtained. 9 when the thickness is 50 μm or more
Thin magnets having a relative density of 0% or more could be produced without cracking or chipping. Next, the molded samples of these thin bulk materials (thicknesses of 1 mm, 0.5 mm, and 0.3 mm, each having an area of 162
FIG. 30 shows the result of X-ray diffraction of the sample (mm ² ).
FIG. 30 shows that bcc- (Fe, Co) and (Fe, Co) ₃
Diffraction lines corresponding to the B or (Fe, Co) ₁₄ Pr ₂ B ₁ phase are observed, and no clear change in the structure due to thinning is observed. Therefore, it has been found that even if the hard magnetic alloy compact of the present invention is made thin, good magnetic properties can be obtained.

Next, FIGS. 31 to 33 show one embodiment of a pen recorder provided with a linear motor to which the thin hard magnetic alloy compact according to the present invention is applied. The pen recorder 20 of this example
Is a thin cassette type, and a mover 23 is slidably provided in the direction of arrow a along a rail-shaped center yoke member 22 erected at the center of a thin box-shaped chassis 21. A pen holder 24 is provided at the distal end, a thin hard magnetic compact (permanent magnet plate) 26 and a bottom yoke member 27 are provided below the center yoke member 22, and a side yoke 28 is provided at an end of the center yoke member 22. ing. A thin hard magnetic compact having a thickness of about 0.5 to 1.3 mm and excellent in hard magnetic properties is required for the thin hard magnetic compact of this configuration. By using the thin hard magnetic compact of the present invention, it is possible to provide a linear motor that exerts a strong driving force even when it is small and thin.

[0083]

As described above, according to the present invention, the element R composed of one or more of the rare earth elements contains 3 to 20 atomic%.
And an alloy made of an Fe-based or FeCo-based alloy containing 2 to 20 atomic% of B, and having a structure in which a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by rapid cooling or an amorphous phase. A soft magnetic phase or a quasi-hard magnetic phase and a mixed phase of a hard magnetic phase are formed in a structure in which a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by crystallization or grain growth under stress, Alternatively, an average crystal grain size of 100 nm
The following fine crystalline phase is precipitated and the mixed phase is formed to form a hard magnetic alloy compact having exchange coupling characteristics in which the fine soft magnetic phase and the hard magnetic phase are combined, and the Fe concentration is reduced. Since it is higher than a conventional rare earth magnet, a strong permanent magnet molded body having high remanence magnetization, squareness ratio and coercive force can be obtained. Further, according to the present invention, since the content of the rare earth element is smaller than that of the conventional rare earth magnet, and excellent hard magnetic properties are obtained, it is possible to obtain a high performance permanent magnet molded article at a relatively low production cost. it can. Further, according to the present invention, the alloy is crystallized or grown under stress to form a soft magnetic phase or a quasi-hard magnetic phase in the structure of the alloy, and a mixed phase state of a hard magnetic phase is formed. Anisotropy is imparted to the hard magnetic phase, which increases uniaxial anisotropy,
High remanence is obtained.

Further, in the present invention, the above alloys are pressed together under pressure to be integrated, so that they are physically more rigid than conventional bonded magnets in which magnetic powder is bound using a binder. In addition, a small permanent magnet having strong hard magnetism can be obtained. Further, in the present invention, the alloy crystallized or grown under stress has a strength of 400 to 10%.
Heat treated at 00 ° C to give an average crystal grain size of 100n in the structure
m or less as a main phase to form a bcc (body-centered cubic structure) -Fe phase or a bcc-Fe
The coercive force of a soft magnetic phase having a Co phase precipitated of 1 kOe or less and the coercive force of a single substance of a Fe ₁₄ R ₂ B phase (where R represents one or more of the rare earth elements) is precipitated. 1k
A hard magnetic phase of Oe or more can be generated, and a permanent magnet having extremely strong hard magnetic properties can be obtained.

In the present invention, an alloy having a structure in which an amorphous phase or a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by quenching is crystallized or grain-grown under stress and consolidated. As a result, a dense hard magnetic alloy compact having a higher relative density than that obtained by sintering after crystallization can be obtained. Further, in the present invention, an alloy containing an amorphous phase that expresses hard magnetism when crystallized is solidified and formed by utilizing a softening phenomenon that occurs during a crystallization reaction, thereby eliminating the need for a binder. A hard magnetic alloy compact having excellent hard magnetic properties and capable of being formed into various shapes can be obtained.

[0086] In the present invention, a composition formula T _x M _y R
represented by _z B _w, where T is, Fe, Co, 1 or more elements of Ni, 1 or more elements of R is a rare earth element, M is Zr, Nb, Ta, Hf, Ti, V, Mo, W
One or more of the above elements, B represents boron, and x, y, z, and w indicating the composition ratio are in atomic%, and 50 ≦ x, 0 ≦ y ≦ 15, 3
By using an alloy represented by ≦ z ≦ 20 and 2 ≦ w ≦ 20, a permanent magnet having strong hard magnetic properties can be obtained.
Furthermore, in the present invention, T _x M _y R _z is represented by B _w E _v, where T is, Fe, Co, 1 or more elements of Ni,
R is one or more of the rare earth elements, M is Zr,
One or more elements of Nb, Ta, Hf, Ti, V, Mo, W, B represents boron, E represents Cr, Al, Pt,
Ru, Rh, Pd, Os, Ir, Cu, Ag, Au, G
x, y, z, w, and v, which represent one or more elements of a and Ge and indicate the composition ratio, are atomic%, and 50 ≦ x, 0 ≦ y ≦
By using an alloy represented by 15, 3 ≦ z ≦ 20, 2 ≦ w ≦ 20, and 0 ≦ v ≦ 10, a permanent magnet having strong hard magnetic properties can be obtained. In the present invention, at least Nd or Pr is contained in the rare earth element R contained in the alloy.
In this case, a permanent magnet having extremely strong hard magnetic properties can be obtained.

Further, in the present invention, by setting the relative density of the compact obtained by compacting the above alloy to be 90% or more, the compact becomes a compact and robust compact. can get. Further, in the present invention, by setting the residual magnetization after consolidation to 100 emu / g or more and / or the ratio of the residual magnetization to the saturation magnetization to 0.6 or more, extremely strong hard magnetic characteristics can be obtained. And a permanent magnet having the same.
Accordingly, the compacted hard magnetic alloy obtained by the present invention is useful as a permanent magnet member used for various devices such as motors, actuators, and speakers, and can be manufactured at low cost. It is an excellent thing that can be reduced.

Further, if the thin hard magnetic alloy compact provided by the present invention can be easily obtained even with a thin hard magnetic property, it can be used as a magnet for a drive device of a thin linear motor or the like. It is suitable and can be used for a wide range of uses at a lower cost than an Sm-Co alloy which has been difficult to process thinly. Further, having the composition of the present invention, if it is a thin hard magnetic alloy compacted body manufactured by compaction, without performing machining such as polishing that was performed on the conventional Sm-Co magnet, thin, large area Since such a magnet is obtained, it is hard to crack as a thin magnet and can be applied to a wide range of uses as a thin magnet.

[Brief description of the drawings]

FIG. 1 is a cross-sectional view showing a main structure of an example of a spark plasma sintering apparatus used to carry out a method for manufacturing a hard magnetic alloy compact according to the present invention.

FIG. 2 is a diagram showing an example of a pulse current waveform applied to an alloy powder in the discharge plasma sintering apparatus shown in FIG.

FIG. 3 is a front view showing the overall configuration of an example of a spark plasma sintering apparatus used for carrying out the method for producing a hard magnetic alloy compact of the present invention.

FIG. 4 is a perspective view for explaining a direction in which a sintering pressure is applied when a compact is manufactured.

FIG. 5 is a graph showing the results of measuring a TMA curve (a) and a DSC curve (b) of a sample of an amorphous alloy according to the present invention and a sample of a crystalline alloy containing no amorphous.

FIG. 6 is a micrograph of the structure of a compact sample obtained when the sintering temperature of the sample of the amorphous alloy according to the present invention is changed in the order of (a), (b), and (c). FIG.

FIG. 7 shows a sample of an amorphous alloy according to the present invention in which a sintering temperature is changed in the order of (a), (b), (c), and (d) to obtain a consolidated sample. It is a graph which shows the pattern by X-ray diffraction.

FIG. 8 is a graph showing the relationship between the sintering temperature and the density of the compact according to the present invention and the relationship between the sintering temperature and the density of the compact of the comparative example.

FIG. 9 is a graph showing the relationship between the sintering pressure and the density of the compact according to the present invention and the relationship between the sintering pressure and the density of the compact of the comparative example.

FIG. 10 is a graph showing an X-ray diffraction result of the amorphous alloy powder according to the present invention.

FIG. 11 is a graph showing an X-ray diffraction result of the amorphous alloy powder according to the present invention.

FIG. 12 shows Fe ₈₈ Nb ₂ Nd ₅ B ₅ and Fe ₈₆ Nb ₂ N
Amorphous powder of d ₇ B ₅ and nano-crystallized Fe ₈₈ Nb ₂ Nd ₅ B ₅
FIG. 3 is a view showing the results of measuring the time (seconds), the temperature and the expansion amount (die displacement amount) of each material, by sintering the powder.

FIG. 13 shows an amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder, an amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder, and a crystalline Fe ₈₈ Nd ₅ Nb.
Molding density of a compact obtained by solidifying and molding ₂ B ₅ powder;
It is a figure showing the relation with sintering temperature.

FIG. 14: Fe ₈₆ Nd ₇ Nb ₂ B ₅ compact and Fe ₈₈
Is a chart showing the X-ray diffraction pattern immediately after sintering of Nd ₅ Nb ₂ B ₅ compacts.

FIG. 15 is a diagram showing a sintering temperature and magnetic properties when amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder and amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder are sintered at a pressure of 636 MPa.

FIG. 16 is a graph showing a relationship between a sintering temperature and magnetic properties when amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder is sintered at a sintering pressure of 636 MPa.

FIG. 17 is a diagram showing the relationship between sintering pressure and magnetic properties when amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder is sintered at a sintering temperature of 600 ° C.

FIG. 18 is a diagram showing the relationship between the magnetic density and the molding density when solidifying and molding amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder.

FIG. 19 shows a Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact and a Fe ₈₆ Nd ₇ solidified at a pressure of 636 MPa and a sintering temperature of 873 K.
For nb ₂ B ₅ compacts is a graph showing the X-ray diffraction results after heat treatment of 750 ° C..

FIG. 20 shows an Fe ₈₈ Nd ₅ Nb ₂ B ₅ compact and a Fe ₈₆ Nd ₇ solidified at a pressure of 636 MPa and a sintering temperature of 600 ° C.
For nb ₂ B ₅ compacts is a diagram showing a magnetization curve after the thermal processing of 750 ° C..

FIG. 21 shows that amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder (Example) and crystalline Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder (Comparative Example) are each at a pressure of 636 MPa and a heating rate of 1.8 K / Seconds,
It is a figure which shows the magnetization curve after heat processing of 750 degreeC about the consolidation body solidified at sintering temperature 600 degreeC and holding time 480 second (8 minutes).

FIG. 22 shows that amorphous Fe ₈₈ Nd ₅ Nb ₂ B ₅ powder
It is a graph which shows the magnetic characteristic after heat processing at 627-827 degreeC about the compact obtained by solidification molding at 36 MPa and sintering temperature 500-600 degreeC.

FIG. 23 shows the pressure of amorphous Fe ₈₆ Nd ₇ Nb ₂ B ₅ powder
It is a graph which shows the magnetic characteristic after heat processing at 627-827 degreeC about the compact obtained by solidifying at 36 MPa and sintering temperature of 500-600 degreeC.

FIG. 24 shows the pressure of amorphous Fe ₉₀ Nb ₂ Nd ₅ B ₅ powder
It is a figure which shows the demagnetization curve of the magnetization of the Example sample heat-processed at 36 MPa and the sintering temperature of 600 degreeC for 8 minutes.

FIG. 25 shows the pressure of amorphous Fe ₈₉ Nb ₂ Nd ₄ B ₅ powder
It is a figure which shows the demagnetization curve of the magnetization of the Example sample heat-processed at 36 MPa and the sintering temperature of 600 degreeC for 8 minutes.

FIG. 26 is a diagram showing a demagnetization curve of magnetization of a sample of an example in which amorphous Fe ₇₆ Co ₁₀ Nb ₂ Nd ₇ B ₅ powder was heat-treated at a pressure of 636 MPa and a sintering temperature of 600 ° C. for 8 minutes.

FIG. 27 is a diagram showing a demagnetization curve of magnetization of a sample of an example in which amorphous Fe ₈₄ Nb ₂ Nd ₇ B ₅ Si ₂ powder was heat-treated at a pressure of 636 MPa and a sintering temperature of 600 ° C. for 8 minutes.

FIG. 28: The maximum magnetic energy product (BH) _max parallel of the sample measured in the pressing direction during molding (Z direction) was measured in a direction perpendicular to the pressing direction during molding (X or Y direction) (BH) _max. It is a figure which shows the dependency between the value divided by vertial and the sintering temperature (Ts).

FIG. 29 is a diagram showing the thickness dependence of the molding density (relative density) of a compact having a composition of Fe ₆₆ Co ₂₀ Nb ₂ Pr ₇ B ₅ .

FIG. 30 shows the thickness of a compact having a composition of Fe ₆₆ Co ₂₀ Nb ₂ Pr ₇ B _{5 of} 1 mm, 0.5 mm, 0.3 mm (area 162).
FIG. ² is an X-ray diffraction diagram of a sample of mm ² ).

FIG. 31 is a diagram showing an example of a pen recorder provided with a thin hard magnetic alloy compact.

FIG. 32 is a bottom view showing a part of the pen recorder shown in FIG. 28.

FIG. 33 is a side view showing a state in which a thin hard magnetic alloy compact applied to the pen recorder shown in FIG. 28 is taken out.

[Explanation of symbols]

A: Spark plasma sintering apparatus, 1: die, 2: upper punch, 3: lower punch, 4, 5: base, 6: alloy powder, 7: Thermocouple, 8: outer frame die, 11: upper base, 12: lower base, 13: chamber.

 ──────────────────────────────────────────────────続 き Continuing on the front page (72) Inventor Akihiro Makino 1-7 Yukiya Otsukacho, Ota-ku, Tokyo Alps Electric Co., Ltd. (72) Inventor Takashi Hatanai 1-7 Yukitani-Otsukacho, Ota-ku, Tokyo Alps Electric Co., Ltd. (72) Inventor Yutaka Yamamoto 1-7 Yukitani Otsuka-cho, Ota-ku, Tokyo Alps Electric Co., Ltd. (72) Inventor Akihisa Inoue 35 Motokawa, Kawauchi, Aoba-ku, Sendai City, Miyagi Prefecture Kawauchi House 11-806

Claims (31)

[Claims] 1. An Fe-based or FeCo-based alloy containing 4 to 20 atomic% of an element R composed of at least one rare earth element and 2 to 20 atomic% of B, and having an average crystal grain size of 100 nm by rapid cooling. An alloy having a structure in which the following microcrystalline phase is precipitated is crystallized or grain-grown under stress, and a soft magnetic phase or a quasi-hard magnetic phase in the structure,
A mixed phase with the hard magnetic phase is formed and anisotropy is given to the hard magnetic phase, so that the coercive force is 1 kOe.
A compact of a hard magnetic alloy characterized by the above. 2. An Fe-based or FeCo-based alloy containing 3 to 20 atomic% of an element R composed of at least one of the rare earth elements and 2 to 20 atomic% of B, and rapidly cooling to form an amorphous phase. The amorphous phase is crystallized under stress in the alloy having a structure including the alloy, and the average crystal grain size is 100 n in the structure.
m or less, and a soft magnetic phase or a quasi-hard magnetic phase is mixed with a hard magnetic phase to form a mixed phase, and the hard magnetic phase is provided with anisotropy, and has a coercive force. Is 1 kOe or more. 3. The hard magnetic alloy compact according to claim 1, wherein the alloy crystallized or grown under stress is subjected to a heat treatment at 400 to 1000 ° C. so that the structure has an average crystal grain size of 100 nm or less. A compact of a hard magnetic alloy, wherein a fine crystalline phase of the above is precipitated as a main phase. 4. The hard magnetic alloy compact according to claim 1, wherein the bcc (body-centered cubic structure) -Fe phase or the bcc-FeCo phase and a Fe-B compound containing a solid solution element. , The coercive force at which at least the amorphous phase is precipitated is 1
a soft or quasi-hard magnetic phase of kOe or less;
Each of the hard magnetic phases having a coercive force of 1 kOe or more in which at least a simple substance of the e ₁₄ R ₂ B phase (where R represents one or more elements among rare earth elements) is precipitated, is 10 vol.
A hard magnetic alloy compact comprising at least (volume)%. 5. A structure in which an amorphous phase or a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by quenching in the hard magnetic alloy compact according to any one of claims 1 to 4. A hard magnetic alloy compact, wherein the alloy is crystallized or grown under stress and compacted. 6. The hard magnetic alloy compact according to claim 1, wherein the alloy containing an amorphous phase and exhibiting hard magnetism when crystallized is softened during a crystallization reaction. A hard magnetic alloy compact formed by solidification utilizing a phenomenon. 7. The compacted hard magnetic alloy according to claim 1, wherein the alloy is heated under stress. 8. The hard magnetic alloy compact according to claim 1, wherein an alloy represented by the following composition formula is used. T _x M _y R _z B _w However, T is, Fe, Co, at least one element of Ni, R is one or more elements of the rare earth elements, M is
One or more elements of Zr, Nb, Ta, Hf, Ti, V, Mo, W, and B represents boron,
y, z, w are atomic%, 50 ≦ x, 0 ≦ y ≦ 15, 3 ≦ z ≦ 2
0, 2 ≦ w ≦ 20. 9. x, y, z, w indicating the composition ratio in the composition formula
Is atomic%, 80 ≦ x ≦ 93, 0.5 ≦ y ≦ 5, 3 ≦ z ≦ 1
The hard magnetic alloy compact according to claim 8, wherein 0 and 3≤w≤7. 10. x, y, which indicate a composition ratio in the composition formula,
z and w are atomic%, 86 ≦ x ≦ 93, 0.5 ≦ y ≦ 3, 3 ≦ z
The hard magnetic alloy compact according to claim 8, wherein ≤7 and 3≤w≤5. 11. The hard magnetic alloy compact according to claim 1, which has the following composition formula. _{T x M y R z B w} E v However, T is, Fe, Co, at least one element of Ni, R is one or more elements of the rare earth elements, M is
One or more elements of Zr, Nb, Ta, Hf, Ti, V, Mo, W, B represents boron, E represents Cr, Al,
Pt, Ru, Rh, Pd, Os, Ir, Cu, Ag, A
x, y, z, w, and v, which represent at least one element of u, Ga, and Ge, and indicate the composition ratio are atomic%, and 50 ≦ x,
0 ≦ y ≦ 15, 3 ≦ z ≦ 20, 2 ≦ w ≦ 20, 0 ≦ v ≦ 10
It is. 12. x, y, which indicate a composition ratio in the composition formula,
z, w, and v are atomic%, 80 ≦ x ≦ 93, 0.5 ≦ y ≦ 5, 3
The hard magnetic alloy compact according to claim 11, wherein ≤ z ≤ 10, 3 ≤ w ≤ 7, and 0 ≤ v ≤ 5. 13. x, y, indicating the composition ratio in the composition formula
z, w, and v are atomic%, and 86 ≦ x ≦ 93, 0.5 ≦ y ≦ 3,
The hard magnetic alloy compact according to claim 11, wherein ≤z≤7, 3≤w≤5, and 0.1≤v≤5. 14. The hard magnetic alloy compact according to claim 8, wherein Si is replaced with T by 0.5 to 0.5%.
A hard magnetic alloy compact comprising 5 atomic% added. 15. The hard magnetic alloy according to claim 1, wherein the rare earth element R contained in the alloy contains at least Nd or Pr. Consolidation. 16. The hard magnetic alloy according to claim 1, wherein a relative density of the compact obtained by consolidating the alloy is 90% or more. Consolidation. 17. The hard magnetic alloy compact according to claim 1, wherein the compact has a residual magnetization of 90 emu / g or more. 18. The hard magnetic alloy compact according to claim 1, wherein the ratio of the residual magnetization (Ir) to the saturation magnetization (Is) is 0.6 or more. Hard magnetic alloy compact. 19. An average crystal of the alloy by rapidly cooling an Fe-based or FeCo-based alloy containing 3 to 20 atomic% of an element R composed of one or more rare earth elements and 2 to 20 atomic% of B. After forming a structure in which a fine crystalline phase having a grain size of 100 nm or less is precipitated, the alloy is crystallized or grown under stress to form a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase in the structure. A method for producing a hard magnetic alloy compact, comprising at least a step of forming a state and imparting anisotropy to the hard magnetic phase. 20. A quenching of an Fe-based or FeCo-based alloy containing 3 to 20 atomic% of an element R composed of at least one of the rare earth elements and 2 to 20 atomic% of B to make the alloy amorphous After forming a structure composed of a crystalline phase, the alloy is crystallized from the amorphous phase under stress to precipitate a fine crystalline phase having an average crystal grain size of 100 nm or less in the structure, and a soft magnetic phase or a quasi-hard phase. A method for producing a hard magnetic alloy compact, comprising at least a step of forming a mixed phase of a magnetic phase and a hard magnetic phase and imparting anisotropy to the hard magnetic phase. 21. The method for manufacturing a compacted hard magnetic alloy according to claim 19, wherein the alloy is crystallized or grown under stress and then heat-treated at 400 to 1000 ° C. to form a structure in the structure. Average crystal grain size 100n
A method for producing a hard magnetic alloy compact comprising precipitating a fine crystalline phase of m or less as a main phase. 22. The method for producing a hard magnetic alloy compact according to claim 19, wherein the alloy is quenched to form an amorphous phase or a fine crystalline phase having an average crystal grain size of 100 nm or less. A method for producing a compacted hard magnetic alloy, comprising, after forming a precipitated structure, crystallizing or growing the alloy under stress and consolidating the alloy. 23. The method for producing a hard magnetic alloy compact according to claim 19, wherein an alloy containing an amorphous phase and exhibiting hard magnetism when crystallized occurs during a crystallization reaction. A method for producing a compacted hard magnetic alloy, which comprises solidifying and molding using a softening phenomenon. 24. The method of manufacturing a hard magnetic alloy compact according to claim 19, wherein the alloy is heated under a stress. 25. The method for producing a hard magnetic alloy compact according to any one of claims 19 to 14, wherein a compact having a relative density of 90% or more is obtained by compacting the alloy. A method of manufacturing a hard magnetic alloy compact. 26. An Fe-based or FeCo-based alloy containing 3 to 20 atomic% of an element R composed of one or more rare earth elements and 2 to 20 atomic% of B, and has a thickness h and an area S.
Wherein the ratio h / S satisfies the relationship of 0.01 ≦ h / S ^0.5 ≦ 0.3, and the relative density is 90% or more. 27. An Fe-based or FeCo-based alloy containing 3 to 20 atomic% of element R consisting of at least one of the rare earth elements and 2 to 20 atomic% of B, and having a thickness of 50 μm.
A thin hard magnetic alloy compact having a relative density of at least 2 mm and a relative density of at least 90%. 28. An alloy having a structure in which a fine crystalline phase having an average crystal grain size of 100 nm or less is precipitated by quenching, is crystallized or grain-grown under stress, and a soft magnetic phase or a quasi-hard magnetic phase is formed in the structure. The thin hard magnetic alloy compact according to claim 26 or 27, wherein a mixed phase of a hard magnetic phase and a hard magnetic phase is formed. 29. An alloy having a structure containing an amorphous phase by quenching, the amorphous phase is crystallized under stress,
A microcrystalline phase having an average crystal grain size of 100 nm or less is precipitated in the structure, and a mixed phase of a soft magnetic phase or a quasi-hard magnetic phase and a hard magnetic phase is formed. 23. The thin hard magnetic alloy compact according to 21 or 22. 30. The hard magnetic alloy compact according to claim 1, having a thickness of 50 μm or more.
mm and a relative density of 90% or more. 31. A hard magnetic alloy compacted body according to claim 19, having a thickness of 50 μm.
A thin hard magnetic alloy compact having a relative density of at least 2 mm and a relative density of at least 90%.