KR20080106211A - R-t-b type alloy and production method thereof, fine powder for r-t-b type rare earth permanent magnet, and r-t-b type rare earth permanent magnet - Google Patents

Abstract

Disclosed is an R-T-B alloy which can be a raw material for a rare earth permanent magnet having excellent magnetic characteristics. Specifically disclosed is an R-T-B alloy containing at least Dy (wherein R represents at least one of Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb and Lu; T represents a transition metal containing not less than 80% by mass of Fe; and B represents a compound containing not less than 50% by mass of B and not less than 0% by mass but less than 50% by mass of at least one of C and N). This R-T-B alloy has a main phase such as an R2T14B phase for exhibiting magnetism, an R-rich phase wherein R is concentrated when compared with the composition ratio of the alloy as a whole, and a Dy concentrated region formed near the R-rich phase in which region Dy is concentrated when compared with the composition ratio of the alloy as a whole. ® KIPO & WIPO 2009

Description

R-T-K-based alloy and its manufacturing method, R-T-K-based rare earth permanent magnet and R-T-K-based rare earth permanent magnet MAGNET, AND RTB TYPE RARE EARTH PERMANENT MAGNET}

The present invention relates to fine powders for R-T-B based alloys, R-T-B based rare earth permanent magnets, and R-T-B based rare earth permanent magnets. Specifically, the present invention relates to a fine powder for R-T-B based rare earth permanent magnets capable of providing an R-T-B based alloy and an R-T-B based rare earth permanent magnet having excellent coercive force.

R-T-B magnets are used in hard disks (HD), magnetic resonance imaging (MRI) and various motors due to their high performance characteristics. In recent years, in addition to improving the heat resistance of R-T-B magnets, the demand for energy saving has increased, so that the utilization rate of a motor including a vehicle motor has increased.

Since R-T-B magnets have Nd, Fe, and B as main components, these types of magnets are collectively referred to as Nd-Fe-B or R-T-B magnets. In RTB-based magnets, R is mainly substituted with a portion of Nd by other rare earth elements such as Pr, Dy, and Tb, T is substituted with a portion of Fe by another transition metal such as Co, Ni, and B is boron It may be partially substituted with C or N as.

RTB-based alloys that can be used in RTB-based magnets are alloys based on a magnetic R ₂ T ₁₄ B phase which contributes to the magnetization action and coexist with a non-magnetic, low-melting R-rich phase enriched with rare earth elements. . Since this RTB-based alloy is an active metal, it is generally melted or cast in a vacuum or inert gas. Usually, the sintered magnet is manufactured from the cast RTB type alloy ingot by the following powder metallurgy process. The alloy ingot is pulverized into an alloy powder having an average particle diameter of about 5 μm (d50: measured by a laser-diffraction particle size analyzer), press-molded in a magnetic field, and sintered at a high temperature of about 1,000 to 1,100 ° C. in a sintering furnace, Thereafter, if necessary, heat treatment and machining are performed, and plating is performed to improve corrosion resistance, thereby completing the manufacture of the sintered magnet.

In R-T-B based sintered magnets, the R-rich phase plays the following important role.

1) Due to the low melting point, it forms a liquid phase during sintering, thereby contributing to the densification of the magnet and consequently to the improvement of magnetization.

2) Eliminates non-uniformity in the grain boundary, thereby reducing nucleation sites in the inertia zone and increasing coercive force.

3) The coercive force is increased by magnetically insulating the columnar column.

Thus, if the R-rich phase in the molded magnet is poorly dispersed, local failure of sintering or deterioration of magnetism is caused. Therefore, it is important that the R-rich phase is uniformly dispersed in the molded magnet. The distribution of the R-rich phase in the R-T-B-based sintered magnet is greatly influenced by the structure of the raw material of the R-T-B-based alloy.

Another problem encountered in casting R-T-B based alloys is the production of α-Fe in the cast alloy. α-Fe is deformable and remains in the mill without being crushed, which not only lowers the crushing efficiency at the time of crushing the alloy but also affects the composition variation and particle size distribution before and after crushing. If α-Fe remains in the magnet even after sintering, the magnetic properties of the magnet are deteriorated. Therefore, in the past, the homogenization treatment was performed on the alloy for a long time at high temperature in order to remove α-Fe. However, since α-Fe exists as a trapezoid nucleus, its removal requires a long time solid phase diffusion. In the case of ingots having a thickness of several centimeters, if the rare earth content is 33% or less, removal of α-Fe is virtually impossible.

In order to solve the problem in which α-Fe is produced in an RTB-based alloy, a strip cast method (simply referred to as "SC method") for casting an alloy ingot at a higher cooling rate has been developed and used. SC method is a method of solidifying an alloy through quenching, and melts an alloy molten metal on the copper roll in which the inside is water-cooled, and produces the flake of about 0.1-1 mm. In the SC method, since the molten alloy is supercooled to a temperature at which the main R ₂ T ₁₄ B phase is formed or to a lower temperature, the R ₂ T ₁₄ B phase can be directly generated from the molten alloy and the formation of α-Fe can be suppressed. Can be. Further, in the SC method, fine microstructures are generated in the alloy, and therefore an alloy having a fine structure that enables fine dispersion of the R-rich phase can be produced. The R-rich phase reacts with hydrogen in a hydrogen atmosphere to expand and become a weak hydride. Using this property, fine cracking suitable for the degree of dispersion of the R-rich phase can be introduced. When the alloy is pulverized through this hydrogenation step, many fine cracks are generated by breaking the hydrogenation trigger of the alloy, so that very good crushability is obtained. Therefore, the internal R-rich phase in the alloy produced by the SC method is finely dispersed, which leads to good dispersibility of the R-rich phase even in the magnet after grinding and sintering, thereby improving the magnetic properties of the magnet (e.g., , Patent Document 1).

In addition, the alloy flakes produced by the SC method are also excellent in the homogeneity of the microstructure. The homogeneity of the microstructures can be compared in terms of crystal grain size or dispersion state of the R-rich phase. In the case of alloy flakes produced by the SC method, chill crystals are sometimes generated on the casting roll side of the alloy flakes (hereinafter referred to as the "mold surface side"), but are appropriately obtained by solidification through quenching as a whole. Fine homogeneous tissue can be obtained. As described above, in the RTB-based alloy produced by the SC method, since the R-rich phase is finely dispersed and the formation of α-Fe is also suppressed, this RTB-based alloy has an excellent microstructure for the production of sintered magnets. have.

The distribution of Dy, which contributes to the improvement of the coercive force, has a great influence on the magnetic properties, in particular the relationship between the coercive force and the elemental distribution in the microstructure of the magnet. For example, the coercive force has already been reported to be high when Dy is distributed near the grain boundary phase (see, for example, Patent Document 2).

More specifically, the coercive force is also reported to be high when Dy is present in the columnar phase (see, for example, Patent Document 3 and Non-Patent Document 1).

In addition, since there is a definite relationship between the magnet properties and the method for producing an alloy, a method for producing an alloy has been advanced along with improvement of the magnet properties. For example, a method of controlling the microstructure (for example, see Patent Document 4) and a method of controlling the microstructure by processing the surface of the casting roll to a predetermined roughness (see, for example, Patent Documents 5 and 6). Known.

[Patent Document 1] Japanese Unexamined Patent Application, First Publication Hei 5-222488

[Patent Document 2] Japanese Unexamined Patent Application, First Publication Hei 5-21219

[Patent Document 3] WO 2003/001541

[Patent Document 4] WO 2005/031023

[Patent Document 5] Japanese Unexamined Patent Application, First Publication No. 2003-188006

[Patent Document 6] Japanese Unexamined Patent Application, First Publication No. 2004-43291

[Non-Patent Document 1] Journal of Japanese Powder and Powder Metallurgy Institute, March 2005, Vol. 52, No. 3, Printing, pages 158-163, Hiroyuki Tomizawa.

However, in recent years, since R-T-B system rare earth permanent magnets having higher performance are desired, there is an increasing demand for further improvement in the magnetic properties of R-T-B system rare earth permanent magnets such as coercive force.

The present invention has been made in view of the above situation, and an object of the present invention is to provide an R-T-B-based alloy as a raw material for a rare earth permanent magnet having excellent magnetic properties.

Still another object of the present invention is to provide a fine powder for an R-T-B rare earth permanent magnet manufactured from the R-T-B base alloy, and an R-T-B rare earth permanent magnet.

The inventors have closely observed the structure of RTB-based alloys containing Dy used to produce RTB-based rare earth permanent magnets in order to investigate the relationship between tissue state and magnetic properties. In addition, the present inventors have found that when the RTB-based alloy containing Dy includes a main phase formed from the R ₂ T ₁₄ B phase and a R-rich phase in which R is concentrated, the Dy-rich region in which Dy is enriched, It was confirmed that RTB-based rare earth permanent magnets obtained by molding / sintering fine powders prepared from flakes of RTB-based alloys would have excellent magnetic properties such as coercive force. The present invention has been completed based on these findings.

That is, the present invention provides the following.

(1) an RTB-based alloy that is used as a rare earth permanent magnet and contains at least Dy (where R is Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, At least one selected from Er, Tm, Yb, and Lu, T is a transition metal containing 80% by mass or more of Fe, B is 50% by mass or more of boron, and 0 to 1 or more elements of C and N. It is contained within a range of less than 50% by mass.), A main phase such as the R ₂ T ₁₄ B phase exhibiting magnetic properties, an R-rich phase in which R is relatively concentrated relative to the total alloy composition ratio, and an R-rich phase An RTB-based alloy formed and comprising a Dy-rich region in which Dy is relatively concentrated compared to the above-described composition ratio.

(2) The R-T-B based alloy according to the above (1), wherein the Dy concentration is lower in the columnar phase than in the Dy-rich region and lower in the R-rich phase than in the columnar phase.

(3) The R-T-B-based alloy according to the above (1) or (2), wherein the alloy is a flake having an average thickness of 0.1 to 1 mm produced by the strip casting method.

(4) A method for producing the RTB-based alloy according to any one of the above (1) to (3), the step of producing a flake having an average thickness of 0.1 to 1 mm, and an average of 10 g / sec or more per 1 cm width Method for producing an RTB-based alloy comprising the step of supplying the molten alloy to the cooling roll at a rate.

(5) The method for producing an R-T-B-based alloy according to the above (4), wherein the R-T-B-based alloy flakes discharged from the cooling roll are held for at least 30 seconds at a temperature of 600 to 900 ° C.

(6) An RTB system produced from the RTB alloy described in any one of (1) to (3) above or from an RTB alloy prepared by the method for producing an RTB alloy described in (4) or (5) above. Fine powder for rare earth permanent magnets.

(7) The R-T-B rare earth permanent magnet manufactured from the fine powder for R-T-B rare earth permanent magnets as described in said (6).

The R-T-B based alloy of the present invention is formed near the R-rich and has a Dy-rich region in which Dy is concentrated relative to the total composition ratio. Therefore, a rare earth permanent magnet having high coercive force and excellent magnetic properties can be realized.

In addition, the fine powder for the RTB rare earth permanent magnet of the present invention and the RTB rare earth permanent magnet are manufactured from either one of the RTB alloy produced by the method of the present invention for producing the RTB alloy or the RTB alloy of the present invention. It has high coercivity and excellent magnetic properties.

1 is a photograph showing an example of the R-T-B-based alloy according to the present invention. This photograph was taken at the time of observing the cross section of the R-T-B type alloy flakes with a scanning electron microscope (SEM).

FIG. 2 is an electron image of the R-T-B based alloy of FIG. 1. FIG.

FIG. 3 is an X-ray image of Fe in the region corresponding to that shown in FIG.

FIG. 4 is an X-ray image of Nd in the region corresponding to that shown in FIG.

FIG. 5 is an X-ray image of Dy in an area corresponding to the area shown in FIG.

FIG. 6 is an X-ray image of Ga in a region corresponding to that shown in FIG.

FIG. 7 is an electron image of the R-T-B based alloy shown in FIG. 1.

FIG. 8 is an X-ray image of Dy in an area corresponding to the area shown in FIG.

FIG. 9 is an X-ray image of Fe in a region corresponding to that shown in FIG.

FIG. 10 is an X-ray image of Nd in the region corresponding to that shown in FIG.

Fig. 11 is a schematic front view showing the construction of an apparatus for producing an alloy according to one embodiment of the present invention.

12 is a schematic front view showing a casting apparatus provided in the apparatus for producing an alloy.

13 is a schematic front view showing a heating device provided in the apparatus for producing an alloy.

14 is a schematic side view showing a heating device provided in the apparatus for producing an alloy.

Fig. 15 is a schematic plan view showing the opening-closing stage and the storage vessel (container) provided in the apparatus for producing an alloy.

Fig. 16 is a schematic front view showing the operation of the apparatus for producing an alloy.

17 is a schematic front view showing the operation of the apparatus for producing an alloy.

18 is a schematic front view showing the operation of the apparatus for producing an alloy.

19 is a schematic front view showing the operation of the apparatus for producing an alloy.

20 is a schematic side view showing the operation of the apparatus for producing an alloy.

Fig. 21 is an electron image of the R-T-B type alloy having no Dy-rich region.

FIG. 22 shows an X-ray image of Dy in an area corresponding to the area shown in FIG. 21.

FIG. 23 shows an X-ray image of Fe in the region corresponding to that shown in FIG.

FIG. 24 shows an X-ray image of Nd in the region corresponding to that shown in FIG.

25 is a graph showing the coercive force (Hcj) of the magnets produced in Examples 1, 2 and Comparative Example 1. FIG.

<Explanation of symbols for the main parts of the drawings>

1: Manufacturing apparatus (device for manufacturing alloys)

2: casting device

3: heating device

4: storage vessel

4a: cooling plate

5: container

6: chamber

7: hopper

7a: hopper outlet

21: grinding device

31: heater

31c: opening

33: open-close stage

33a: stage plate

33b: opening and closing system

51: belt conveyor (moving device)

L: molten alloy

N: flakes of casting alloy

1 is a photograph showing an example of the R-T-B-based alloy of the present invention. This photograph was taken at the time of observing the cross section of the R-T-B type alloy flakes with a scanning electron microscope (SEM). In Fig. 1, the left side is the mold surface.

The R-T-B based alloy shown in Fig. 1 is produced by the SC method. This R-T-B-based alloy has a composition by weight of 23% Nd, 9% Dy, 1% B, 1% Co, 0.2% Ga, and the balance Fe. RTB alloy of the present invention, wherein R is at least one selected from Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb, and Lu, and T Is a transition metal containing 80% by mass or more of Fe, B contains 50% by mass or more of boron, and contains at least one element of C and N within a range of 0 to 50% by mass or less). It is not limited to a specific composition, and the alloy may have any composition as long as it is an RTB-based alloy containing at least Dy.

The RTB-based alloy shown in FIG. 1 is composed of an R ₂ T ₁₄ B phase (main phase) and an R-rich phase. In Fig. 1, the R-rich phase is shown in white and the R ₂ T ₁₄ B phase (main phase) is shown in gray. R ₂ T ₁₄ B phase is mainly formed of columnar tablets and partially formed of equiaxed crystals. The average grain size in the short axis direction of the R ₂ T ₁₄ B phase is 10 to 50 µm. In the R ₂ T ₁₄ B phase, linear R-rich phases or particulate or partially broken R-rich phases extending along the major axis direction of the columnar tablet exist in grain boundaries and grain boundaries. The R-rich phase is a non-magnetic phase having a low melting point and enriched in R relative to the total composition. The average distance between R-rich phases is 3-10 μm.

2-6 show the results of elemental distribution analysis (digital mapping) of the R-T-B based alloy shown in FIG. 1 by an electron microscopy analyzer (EPMA) using a wavelength dispersed X-ray spectrometer (WDS).

FIG. 2 is an electronic image of the RTB-based alloy shown in FIG. 1. The R-rich phase is shown in white and the R ₂ T ₁₄ B phase (main phase) is shown in gray.

FIG. 3 is an X-ray image of Fe in the region corresponding to that shown in FIG. It is clear from Figures 2 and 3 that the R-rich phase contains less Fe as compared to the main phase.

FIG. 4 is an X-ray image of Nd in the region corresponding to that shown in FIG. From Figures 2 and 4 it is clear that the R-rich phase contains more Nd compared to the main phase.

FIG. 5 is an X-ray image of Dy in an area corresponding to the area shown in FIG. It is clear from Figures 2 and 5 that the R-rich phase contains less Dy as compared to the main phase.

FIG. 6 is an X-ray image of Ga in a region corresponding to that shown in FIG. It is clear from Figures 2 and 6 that the R-rich phase contains more Ga as compared to the main phase.

7-10 show the results of elemental distribution analysis (digital mapping) using field emission-electron microscopy (FE-EPMA).

FIG. 7 is an electronic image of the RTB-based alloy shown in FIG. 1. The R-rich phase is shown in white and the R ₂ T ₁₄ B phase (main phase) is shown in gray.

FIG. 8 is an X-ray image of Dy in an area corresponding to the area shown in FIG. It is clear from Figs. 7 and 8 that the Dy-rich region in which Dy is relatively concentrated in comparison with the R-rich phase and the main phase is formed closer to the R-rich phase. It is also clear from FIG. 8 that the Dy concentration is lower in the columnar phase and lower in the R-rich phase than in the Dy-rich region.

FIG. 9 is an X-ray image of Fe in a region corresponding to that shown in FIG. It is clear from Figures 7 and 9 that the R-rich phase contains less Fe as compared to the main phase.

FIG. 10 is an X-ray image of Nd in the region corresponding to that shown in FIG. It is clear from Figures 7 and 10 that the R-rich phase contains more Nd compared to the main phase.

(Production method)

The R-T-B-based alloy of the present invention shown in FIG. 1 can be cast, for example, by the SC method using the apparatus for producing an alloy shown in FIG.

[Alloy manufacturing apparatus]

11 is a schematic front view showing the overall configuration of the apparatus for producing an alloy of the present invention.

The apparatus 1 for producing an alloy shown in FIG. 11 (hereinafter referred to as "manufacturing apparatus 1") is usually provided with a casting apparatus 2, a grinding apparatus 21, and a heating apparatus 3. The heating device 3 comprises a heater 31 and a container 5. The container 5 comprises a storage vessel 4 and an opening-closing stage group 32 provided above the storage vessel 4.

The manufacturing apparatus 1 shown in FIG. 11 has a chamber 6. The chamber 6 includes a casting chamber 6a and a temperature-keeping storage chamber 6b provided below the casting chamber 6a and connected to the casting chamber 6a. The casting device 2 is installed in the casting chamber 6a, and the heating device 3 is installed in the temperature-keeping storage chamber 6b. In addition, a gate 6e is provided in the temperature-keeping storage chamber 6b, and the temperature-keeping storage chamber 6b is except when the container 5 is transferred outside the temperature-keeping storage chamber 6b. It is closed by the gate 6e.

In addition, the casting device 2 is provided with a grinding device 21 and a hopper 7 is provided between the casting device 2 and the open-close stage group 32. The hopper 7 leads the cast alloy flakes to the open-close stage group 32.

[Casting device]

12 is a schematic front view showing the casting apparatus 20 provided in the manufacturing apparatus 1.

The casting apparatus 2 shown in Fig. 12 is a cooling roll 22 for casting the molten alloy L into a cast alloy M by a method of quenching the molten alloy using a water cooling system (not shown in the drawing). ; A tundish 23 for supplying the molten alloy L to the cooling roll 22; And a grinding device 21 for grinding the cast alloy M into cast alloy flakes N. As shown in FIG. As shown in Fig. 12, the grinding device 21 includes, for example, a pair of grinding rolls 21a.

[Heater]

FIG. 13 is a schematic front view showing a heating device 3 provided in the manufacturing apparatus 1, FIG. 14 is a schematic side view thereof, and FIG. 15 is a schematic plan view thereof.

As shown in Figs. 13 to 15, the heater 31 included in the heating device 3 has a heater cover 31a and a main body 31b attached below the heater cover 31a. In order to discharge heat generated from the main body 31b in the direction of the container 5 and to prevent heat from the main body 31b from being discharged into the casting chamber 6a, a heater cover 31a is provided. In addition, if the heater cover 31a is provided, this can prevent the main body 31b from being damaged when a part of the molten alloy or the cast alloy falls unexpectedly into the main body.

The heater 31 has the opening part 31c, and the discharge port 7a of the hopper 7 is arrange | positioned at the opening part 31c. As a result, the flakes N of the cast alloy falling from the casting apparatus 2 after passing through the hopper 7 are supplied to the open-close stage group 32 in the container 5 provided under the heater 31. Can be.

11 and 13, the heater 31 is along the longitudinal direction (moving direction of the container 5) of the belt conveyor 51 provided inside the temperature-keeping storage chamber 6b. Is placed.

With this configuration, even if the container 5 moves into the temperature-maintaining storage chamber 6b, the temperature of the flakes N of the cast alloy placed on the open-close stage group 32 in the container 5 is uniform. Can be maintained.

The opening-closing stage group 32 included in the heating device 3 is integrated with the storage vessel 4 to form the container 5. That is, the container 5 shown in FIGS. 13 to 15 is formed of a storage vessel 4 and an opening-closing stage group 32 provided over the storage vessel 4.

The opening-closing stage group 32 has a plurality of opening-closing stages 33 arranged along the moving direction of the container 5. In addition, a guide member 52 is provided around the opening-closing stage group 32 in which the flakes N of the cast alloy falling through the hopper 7 are stored in the temperature-keeping storage chamber ( 6b) to prevent it from being scattered inside;

Each of the opening-closing stages 33 is supplied from the casting apparatus 2 to hold the flakes N of the casting alloy mounted thereon to maintain the temperature by the heater 31 in a predetermined period, and after the temperature holding time The flakes N of the cast alloy are dropped into the storage vessel 4. Each open-close stage 33 has a stage plate 33a and an open-close system 33b that opens or closes the stage plate 33a. Each opening-closing system 33b includes a rotating shaft 33b ₁ attached to one side of the stage plate 33a; And a drive unit (not shown) for rotating the rotation shaft 33b ₁ . Each drive unit can freely rotate the rotation shaft 33b ₁ so that the inclination angle of each stage plate 33a can be individually controlled. The angle of inclination of each stage plate 33a is 90 ° clockwise from 0 ° (position where the stage plate 33a is horizontal (position shown in Fig. 13 by the dashed-dotted line)) (position where the stage plate 33a is almost vertical ( Can be set to any angle within the range of [] shown by the solid line).

Thus, the open-close stage 33 can leave the flakes N of the cast alloy mounted on the stage plate 33a for the predetermined temperature holding time by driving the open-close system 33b, and then By making the inclination angle of the stage plate 33a larger, the flakes N of the cast alloy can be dropped into the storage vessel 4.

In addition, the opening-closing stage 33 can serve as a cover for the storage vessel 4, which prevents the heat of the heater 31 from reaching the storage vessel 4. Prevent the interior from heating up. In addition, a plurality of cooling plates 4a are provided inside the storage vessel 4.

Further, as shown in Figs. 13 and 14, the container 5 is mounted on the belt conveyor 51 (moving device). The belt conveyor 51 causes the container 5 to move left or right in FIG.

[Alloy casting]

16 to 19 are schematic front views illustrating the operation of the apparatus for producing an alloy.

As shown in Fig. 16, the container 5 firstly has the opening-closing stage 33A (which is present at the left edge of the opening-closing stage group 32 in the drawing) with the outlet 7a of the hopper 7. It is moved to where it is located directly below. In addition, all the open-close stages 33 are set to the closed state.

Subsequently, the casting apparatus 2 shown in FIG. 12 is driven to prepare the flakes N of the cast alloy. First, the molten alloy L is prepared in a melting apparatus (not shown in the figure). The temperature of the molten alloy L varies depending on the type of alloy contents, but is controlled within the range of 1,300 ° C to 1,500 ° C. The prepared molten alloy L is transferred to the casting apparatus 2 while being held in the refractory crucible 24. Subsequently, the molten alloy L is supplied from the refractory crucible 24 to the tundish 23 and subsequently supplied from the tundish 23 to the cooling roll 22, where the molten alloy L is solidified and cast. Alloy M is produced. Thereafter, the casting alloy M is moved from the cooling roll 22 to the opposite side of the tundish 23, and two rotating grinding rolls are formed so that the casting alloy M is crushed into the flakes N of the casting alloy. 21a).

The average alloy molten metal feed rate to the cooling roll 22 is 10 g / sec or more per cm of width, preferably 20 g / sec or more, more preferably 25 g / sec or more, even more preferably 100 g per cm of width / sec or less. When the feed rate of the molten alloy L is less than 10 g / sec, the molten alloy L is thinly wetted and does not spread on the cooling roll 22, but instead of the viscosity of the alloy molten L itself or the casting roll 22 Due to the wettability to the surface, shrinkage can lead to variations in alloy quality. On the other hand, when the average alloy molten metal supply rate to the cooling roll 22 exceeds 100 g / sec per 1 cm in width, cooling on the cooling roll 22 is insufficient, resulting in coarsening of fine structures, precipitation of α-Fe, and the like. Can be.

It is preferable that the average cooling rate of the molten alloy on the cooling roll 22 is 100-2,000 degreeC / sec. The average cooling rate of 100 ° C./sec or more is satisfactory to prevent precipitation of α-Fe or tissue coarsening on the R-rich phase. On the other hand, if the average cooling rate is 2,000 ° C./sec or less, the degree of supercooling may be excessive and the cast alloy flakes may be supplied to the heating device 3 at an appropriate temperature. In addition, the cast alloy flakes are not overly cooled and do not require a reheating process. An average cooling rate is calculated | required by dividing the difference between the temperature just before the contact of a cooling roll and a molten alloy, and the temperature at which it removes from a cooling roll by the time during which the molten alloy is in contact with a cooling roll.

The average temperature of the cast alloy M at the time of separation from the cooling roll 22 is a slight difference in the degree of contact between the casting alloy M and the cooling roll 22, a variation in the thickness of the casting alloy M, and the like. Due to this fluctuate slightly. The average temperature of the cast alloy M when detached from the cooling roll 22 is measured, for example, by scanning the alloy surface in the width direction from the beginning to the end of the casting with a radiation thermometer and measuring the temperature. This can be obtained by averaging the measured values.

The average temperature of the cast alloy M when separated from the cooling roll 22 is preferably 100 to 500 ° C. lower than the solidification temperature of the R ₂ T ₁₄ B phase in the equilibrium state of the molten alloy, and 100 to 400 ° C. lower. More preferred. The melting temperature of the R ₂ T ₁₄ B phase was found to be 1,150 ° C. in the Nd-Fe-B ternary system, but Nd was substituted with another rare earth element, Fe was substituted with another transition element, and the kind of optional additional elements and It varies depending on the amount added. If the difference between the average temperature of the cast alloy M when it is separated from the cooling roll 22 and the solidification temperature of the R ₂ T ₁₄ B phase in the equilibrium state of the cast alloy M is less than 100 ° C, this is insufficient. It corresponds to the cooling rate. On the other hand, if the difference exceeds 500 DEG C, the supercooling of the molten alloy becomes too large because of the too high cooling rate.

In addition, the average temperature of the cast alloy M when separated from the cooling roll 22 also fluctuates within the same casting step (tap), and if the fluctuation range is large, this may cause fluctuations in the microstructure or quality. Therefore, it is appropriate that the fluctuation range of the temperature in the tap is narrower than 200 ° C, preferably 100 ° C or less, more preferably 50 ° C or less, even more preferably 20 ° C or less.

The cast alloy flakes N preferably have an average thickness of 0.1 to 1 mm. If the average thickness of the flakes is less than 0.1 mm, the solidification rate may be too high and the R-rich phase may be dispersed too finely. On the other hand, if the average thickness of the flakes exceeds 1 mm, the solidification rate decreases, which may cause a decrease in the dispersibility of the R-rich phase, precipitation of α-Fe, and the like.

Next, as shown in Fig. 16, the thin laminas N of the cast alloy are transferred to the heating apparatus 3 by passing through the hopper 7, and are located directly under the outlet 7a of the hopper 7. Stacked (closed) on the closed stage 33A. During this time, the heater 31 is turned on and the flakes N of the cast alloy are kept at a predetermined temperature or heated by the heater 31 immediately after they are stacked on the opening-closing stage 33A.

The amount of the flakes N of the cast alloy stacked on the opening-closing stage 33A can be appropriately adjusted according to the area of the stage plate 33a. However, since the flakes N of the cast alloy are continuously supplied from the casting apparatus 2, depending on the feed rate, they will overflow from the opening-closing stage 33A after time. For this reason, when the stacking amount of the flakes N of the cast alloy reaches a predetermined value with respect to the opening-closing stage 33A, the container 5 is moved to the left in the figure as shown in FIG. Subsequently, another open-close stage 33B next to the right-opening stage 33A is directly below the outlet 7a of the hopper 7, followed by a flake N of cast alloy stacked on the open-close stage 33B. Is located in. Then, in the same manner, the container 5 is moved in accordance with the preparation of the flakes N of the cast alloy, and the flakes N of the cast alloy are sequentially stacked on the opening-closing stages 33C to 33E.

The flakes N of the cast alloy stacked on each of the opening-closing stages 33A to 33E are maintained at a predetermined temperature or heated by the heater 31. The holding temperature is preferably lower than the temperature (separation temperature) of the flakes N when separated from the cooling roll. Specifically, the holding temperature is preferably in the range from (separation temperature-100 ° C) to the separation temperature, and (separation It is more preferable to exist in the range from the temperature -50 degreeC) to a separation temperature. More specifically, the holding temperature is preferably in the range of 600 ° C to 900 ° C. If the holding temperature is 600 ° C or higher, the coercive force of the R-T-B-based alloy can be sufficiently improved. In addition, when the holding temperature is 900 ° C. or less, precipitation of α-Fe can be prevented, and fine structure such as an R-rich phase can be prevented from being coarsened.

In addition, when the separation temperature is lowered for some reason, by setting the holding temperature higher than the separation temperature, the flakes N of the cast alloy can be heated and kept at a predetermined temperature. It is preferable to exist in 100 degreeC, and, as for a heating range, it is more preferable to exist in 50 degreeC. If the heating range is too wide, the manufacturing efficiency is lowered. Even if the flakes are kept at 1,000 ° C., the coercive force can be improved. However, this temperature harmonizes the microstructure. It is also disadvantageous because the particle distribution or fluidity of the fine powder and the sintering temperature fluctuate when finely ground. Therefore, when they are maintained at 1,000 ° C., it is desirable to consider their impact on subsequent processes.

In addition, the temperature holding time is preferably 30 seconds or more, more preferably from 30 seconds to several hours, and most preferably from 30 seconds to about 30 minutes. If the temperature holding time is 30 seconds or more, the coercive force can be sufficiently improved. That is, the flakes of the cast alloy may be subjected to temperature holding treatment for several hours, but from the viewpoint of manufacturing efficiency, the temperature holding time is preferably 30 seconds or less.

Next, as shown in Fig. 18, in accordance with the preparation of the flakes N of the cast alloy, the container 5 is further moved in the same manner with respect to the remaining opening-closing stages 33F to 33J, thereby casting the cast alloy. Flakes N are successively stacked on each opening-closing stage 33F to 33J. With respect to the flakes N of the cast alloy stacked on the open-close stages 33A to 33D, when the predetermined temperature holding time elapses, each open-close stage is sequentially opened as shown in FIG. By making the flakes fall continuously into the storage vessel 4. When the flakes N of the cast alloy fall into the storage vessel 4, the heat of the heater 31 no longer reaches the flakes N of the cast alloy and thus the temperature holding process is terminated.

As described above with reference to Fig. 17, since the flakes N of the cast alloy are successively mounted on each opening-closing stage, different opening-closing stages are attached to the flakes N of the casting alloy on the opening-closing stage. There is a time difference at the starting point of starting the temperature holding process. Therefore, in order to fix the temperature holding time with respect to the flakes N of the cast alloy on each open-close stage, the flakes N of the cast alloy are stored in the storage vessel by successively switching each open-close stage to an open state. (4) It is preferable to fall continuously into.

The flakes N of the cast alloy dropped into the storage vessel 4 come into contact with the cooling plate 4a, whereby heat is absorbed into the cooling plate 4a, whereby the flakes N of the cast alloy are cooled.

19 and 20 show a state in which all the opening-closing stages 33 are in an open state and the flakes N of the cast alloy are stored in the storage vessel 4. Then, if casting and grinding by the casting apparatus 2 are subsequently performed, the container 5 can be moved to the right in the drawing while all the opening-closing stages 33 are set in the closed state, The flakes N are continuously mounted on each opening-closing stage 33 in accordance with the preparation of the flakes N of the cast alloy. In contrast, when the casting and grinding process by the casting device 2 is finished, all the opening-closing stages 33 are switched to the closed state to prevent the heat of the heater 31 from reaching the storage vessel 4. . Subsequently, the gate 6e of the temperature-holding chamber 6b is opened, and the container 5 is transferred out of the chamber 6 to collect the flakes N of the cast alloy, thereby producing the cast alloy flakes N. To complete.

[Cooling speed]

Next, the cooling rate at the time of manufacturing an R-T-B type alloy is demonstrated.

In the present invention, the cooling rate was controlled to obtain the next cooling rate from the solidification point (about 1,170 ° C) of the main phase which is the temperature immediately after solidification to 600 ° C which is lower than the freezing point (about 700 ° C) of the R-rich phase.

That is, the cooling rate of R-T-B type alloy is set in the range of 100-300 degreeC / sec from 1,000 degreeC to 850 degreeC. If the cooling rate from 1,000 ° C. to 850 ° C. is higher than this range, Dy may not diffuse sufficiently into the columnar phase. On the other hand, if the cooling rate is lower than the above range, Dy diffuses excessively, making it impossible to form the Dy-rich region on the columnar phase.

Moreover, it is preferable that the cooling rate of an R-T-B type alloy is set to the range of 300-2,000 degreeC / sec from the solidification point of a columnar phase to 1,000 degreeC. By setting the cooling rate of the R-T-B-based alloy from the solidification point of the columnar to 1,000 ° C within the above range, the R-T-B-based alloy having the Dy-rich region is obtained with high productivity.

Moreover, it is preferable to temporarily set the cooling rate of R-T-B type alloy to 100 degrees C / sec or less from 850 degreeC to 600 degreeC. By temporarily setting the cooling rate within the above range from 850 ° C to 600 ° C, Dy contained in the R-rich phase can sufficiently diffuse into the adjacent columnar phase. Therefore, an R-T-B based alloy having a Dy-rich region and a higher coercive force can be easily manufactured.

The flakes of the R-T-B-based alloy of the present embodiment and the R-T-B-based alloy are formed close to the R-rich and have a Dy-rich region in which Dy is relatively concentrated compared to the total composition ratio. Therefore, rare earth permanent magnets having high coercive force and excellent magnetic properties can be obtained from them.

In other words, the R-T-B-based alloy of the present embodiment has a higher coercive force, for example, as compared with the R-T-B-based alloy shown in Figs. 21 to 24 having no Dy-rich region.

21-24 show the results of elemental distribution analysis (digital mapping) of R-T-B based alloys having no Dy-rich region as an example using a field emission-electron microscopic analyzer (FE-EPMA). The R-T-B based alloys shown in Figs. 21 to 24 are manufactured by the SC method. This R-T-B-based alloy has a composition by weight of 23% Nd, 9% Dy, 1% B, 1% Co, 0.2% Ga, and the balance Fe.

21 is an electron image of an RTB-based alloy having no Dy-rich region. The R-Li phase is shown in white and the R ₂ T ₁₄ B phase (main phase) is shown in gray.

FIG. 22 is an X-ray image of Dy in an area corresponding to the area shown in FIG. It is clear from Figs. 21 and 22 that this R-T-B based alloy has a higher Dy and a lower Dy-rich region on the R-rich than the main phase compared to the main phase.

FIG. 23 is an X-ray image of Fe in the region corresponding to that shown in FIG. It is apparent from FIGS. 21 and 23 that the R-rich phase contains less Fe as compared to the main phase.

FIG. 24 is an X-ray image of Nd in the region corresponding to that shown in FIG. It is clear from Figs. 21 and 24 that the R-rich phase contains more Nd compared to the main phase.

[Manufacture of R-T-B Type Rare Earth Permanent Magnet]

In order to manufacture the R-T-B based rare earth permanent magnet of the present invention, first, a fine powder for the R-T-B based rare earth permanent magnet is prepared from the R-T-B based alloy of the present invention. The fine powder for RTB-based rare earth permanent magnets of the present invention is obtained by, for example, performing hydrogen cracking of flakes formed of the RTB-based alloy of the present invention by hydrogen absorption and then pulverizing the flakes using a grinder such as a jet mill. . In the hydrogen cracking here, for example, it is preferable that a hydrogen absorption step of holding the flakes in a hydrogen atmosphere under a predetermined pressure is performed in advance.

Subsequently, the obtained fine powder for R-T-B-based rare earth permanent magnets is press-molded, for example, by a molding machine or the like in a transverse magnetic field and sintered in vacuum to obtain an R-T-B-based rare earth permanent magnet.

The fine powder for the R-T-B-based rare earth permanent magnet and the R-T-B-based rare earth permanent magnet of the present embodiment is produced from the R-T-B-based alloy of the present invention. Therefore, they have high coercive force and excellent magnetic properties.

Example 1

Weigh the starting material formulated to have an alloy composition of 23% Nd, 9% Dy, 0.98% B, 1% Co, and 0.2% Ga, balance Fe, and then use a high frequency melting furnace to The molten alloy was prepared by melting in an alumina crucible in an argon gas atmosphere. Then, this alloy molten metal was supplied to the casting apparatus of the manufacturing apparatus shown in Fig. 11 and cast by the SC method. The rotational speed of the cooling roll was 1.3 m / s at the time of casting, the average alloy melt supply rate to the cooling roll was 30 g / sec per cm of width, and the average temperature of the cast alloy ingot was 850 ° C. upon separation from the cooling roll. .

The cooling rate of this alloy was 700 ° C / sec from the solidification point of the columnar to 1,000 ° C, 200 ° C / sec from 1,000 ° C to 850 ° C, and 50 ° C / sec from 850 ° C to 780 ° C. Thereafter, the alloy was held at a temperature of about 780 ° C for 300 seconds on the opening-closing stage using the manufacturing apparatus shown in Fig. 11, and then cooled to 600 ° C or lower at a cooling rate of 0.1 ° C / sec. Example 1 A thin piece of RTB alloy was prepared. At this point, the average thickness of the alloy was 0.3 mm.

Example 2

The molten alloy was prepared using the same apparatus as the same starting material as in Example 1. Next, the obtained molten alloy was cast using the same casting apparatus as in Example 1. The rotational speed of the cooling roll at the time of casting was 0.87 m / s, the average alloy melt supply rate to the cooling roll was 30 g / sec per cm of width, and the average temperature of the casting alloy ingot at the time of separation from the cooling roll was 880 ° C.

The cooling rate from the solidification point of the main phase of this alloy to 1,000 degreeC was 700 degreeC / sec, was 200 degreeC / sec from 1,000 degreeC to 850 degreeC, and was 10 degreeC / sec from 850 degreeC to 780 degreeC. Thereafter, using the manufacturing apparatus shown in Fig. 11 without using the opening-closing stage, the alloy was cooled to 600 ° C or lower at a cooling rate of 0.1 ° C / sec to prepare a flake of the R-T-B-based alloy of Example 2. At this point, the average thickness of the alloy was 0.45 mm.

The flakes of the RTB-based alloys obtained in Examples 1 and 2 were prepared using an electron microscope analyzer (WDS-EPMA) equipped with a wavelength dispersion X-ray spectrometer and a field emission electron microscope analyzer (FE-EPMA). Distribution analysis (digital mapping) (surface analysis) was performed. As a result, it was found that both the flakes of the R-T-B-based alloys obtained in Examples 1 and 2 form a Dy-rich region in which Dy is concentrated in comparison with the R-rich phase and the main phase near the R-rich phase. In addition, in both the flakes of the R-T-B-based alloys obtained in Examples 1 and 2, the Dy concentration was lower in the columnar phase than in the Dy-rich region, and even lower in the R-rich phase.

Comparative Example 1

The molten alloy was prepared using the same apparatus as the same starting material as in Example 1. Subsequently, the obtained molten alloy was cast using the same casting apparatus as in Example 1 to prepare flakes of the R-T-B based alloy of Comparative Example 1. The rotational speed of the cooling roll at the time of casting was 0.65 m / s, the average alloy melt supply rate to the cooling roll was 15 g / sec per cm of width, and the average temperature of the casting alloy ingot at the time of separation from the cooling roll was 700 ° C.

The cooling rate from the solidification point of the main phase of this alloy to 1,000 degreeC was 700 degreeC / sec, 400 degreeC / sec from 1,000 degreeC to 700 degreeC, and 10 degreeC / sec from 700 degreeC to 600 degreeC. Thereafter, the alloy was cooled to 600 ° C. or lower at a cooling rate of 0.1 ° C./sec using the manufacturing apparatus shown in FIG. 11 without using the open-close stage. At this point, the average thickness of the alloy was 0.30 mm.

Elemental distribution analysis (digital mapping) (surface analysis) was performed on the flakes of the R-T-B-based alloy obtained in Comparative Example 1 using WDS-EPMA and FE-EPMA. As a result, it was found that the flakes of the R-T-B based alloy obtained in Example 1 did not form any Dy-rich region in which Dy was concentrated compared with the total composition ratio. One possible reason for this result is that in Comparative Example 1, when the alloy alloy ingot is low when separated from the cooling roll, and the alloy that is too quenched on the cooling roll, the cooling rate of the alloy from 1,000 ° C to 700 ° C is too high. Because of making. Therefore, maybe Dy and Nd were not sufficiently diffused and no concentration gradient was formed.

Next, using the flakes of the R-T-B-based alloy obtained in Example 1, Example 2 and Comparative Example 1 to prepare a magnet as follows.

Hydrogen cracking was first performed on flakes of the R-T-B based alloys obtained in Examples 1, 2, and Comparative Example 1. Hydrogen cracking was performed by the following method. The flakes of the R-T-B based alloy were made to absorb hydrogen in a hydrogen atmosphere of 2 atm, and then heated to 500 ° C. in a vacuum to remove hydrogen. Thereafter, 0.07 mass% of zinc stearate was added thereto and the result was ground by a jet mill using a nitrogen gas stream. The powder obtained by milling had an average particle size of approximately 5.0 μm as measured by laser diffraction.

Next, the obtained powder material was press-molded by the molding machine in the transverse magnetic field in the molding pressure of 0.8t / cm <2> in 100% nitrogen atmosphere, and the powder green compact was obtained. The resulting powder green compact was heated from room temperature in a vacuum of 1.33 × 10 ⁻⁵ hPa, maintained at 500 ° C. for 1 hour, and then maintained at 800 ° C. for 1 hour to remove zinc stearate and remaining hydrogen. Subsequently, the obtained powder green compact was heated to a sintering temperature of 1,030 ° C. and held there for 3 hours to prepare a sintered compact. Thereafter, the obtained sintered body was subsequently heat treated at 800 ° C. and at 530 ° C. in an argon atmosphere for 1 hour. As a result, ten magnets were obtained in both Example 1 and Example 2, and five magnets were obtained in Comparative Example 1.

Magnetic properties of the magnets obtained in Examples 1, 2 and Comparative Example 1 were measured by a DC BH curve tracer. The results are shown in Table 1 and FIG. FIG. 25 is a graph showing the coercive force (Hcj) of the magnets prepared in Examples 1, 2, and Comparative Example 1. FIG. In this graph, the vertical axis represents the coercive force of each embodiment shown in the horizontal. In Fig. 25, symbol o indicates the coercive force of the magnets produced in Examples 1 and 2, and symbol o indicates the coercive force of the magnets produced in Comparative Example 1.

In Table 1, "(BH) _max " represents maximum magnetic energy, "Br" represents residual magnetic flux density, "Hcj" represents coercive force, and "Hk / Hcj" represents hysteresis angulation.

As shown in Table 1 and Figure 25, the magnets obtained in Examples 1 and 2 were compared with the coercive force of the magnets obtained in Comparative Example 1 prepared from an RTB-based alloy in which no Dy-rich region was formed. Had a high coercivity "Hcj". This difference in the coercive forces of the magnets is derived when they are in the alloy state and is caused by the distribution of element concentrations therein that continue to influence even after the alloy has been ground and sintered to produce the magnets. Since there are Dy-rich regions in the alloy of the present invention and they also remain in the grains of the magnet, it is possible that only a small amount of Dy remains on the R-rich which does not contribute effectively to the improvement of the coercive force.

Claims (7)

It is a raw material used for rare earth permanent magnets and is an RTB alloy containing at least Dy (where R is Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm , Yb, and Lu are at least one selected from T, T is a transition metal containing 80% by mass or more of Fe, B contains 50% by mass or more of boron, and 0 to 50% by mass of one or more elements of C and N. It is contained in the following range.), Columnar phases such as the R ₂ T ₁₄ B phase exhibiting magnetic properties, R-rich phase in which R is concentrated compared to the total alloy composition ratio, An R-T-B based alloy formed near a R-rich phase and comprising a Dy-rich region in which Dy is concentrated compared to the composition ratio. The R-T-B based alloy of claim 1, wherein the Dy concentration is lower in the columnar phase than in the Dy-rich region and lower in the R-rich phase than in the column phase. The R-T-B-based alloy according to claim 1 or 2, wherein the alloy is a flake having an average thickness of 0.1 to 1 mm produced by the strip casting method. The method of manufacturing the R-T-B-based alloy according to any one of claims 1 to 3, Preparing a flake having an average thickness of 0.1 to 1 mm, A method for producing an R-T-B alloy comprising the step of supplying molten alloy to the cooling roll at an average speed of 10 g / sec or more per 1 cm in width. The method for producing an R-T-B-based alloy according to claim 4, wherein the R-T-B-based alloy flakes discharged from the cooling roll are maintained for at least 30 seconds at a temperature of 600 to 900 ° C. For RTB-based rare earth permanent magnets produced from the RTB-based alloy according to any one of claims 1 to 3, or from an RTB-based alloy prepared by the method for producing an RTB-based alloy according to claim 4 or 5. Fine powder. An R-T-B rare earth permanent magnet prepared from a fine powder for an R-T-B rare earth permanent magnet according to claim 6.

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