RU2389097C1 - Alloy of r-t-b type and procedure for its production, fine dispersed powder for rare earth permanent magnet r-t-b type and rare earth permanent magnet of r-t-b type - Google Patents

Abstract

FIELD: metallurgy.

SUBSTANCE: alloy R-T-B containing Dy, where R is at least one of Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb and Lu, and T is transition metal containing not less, than 80 wt % Fe; B is composition containing not less, than 50 wt % B and from 0 to less, than 50 wt % at least one from C and N. Alloy has a main phase R₂T₁₄B to facilitate magnet properties, a rich R phase, wherein R is concentrated in comparison with its share in general contents of alloy and a rich Dy region formed near a rich R phase, wherein Dy is concentrated in comparison with its share in a general contents of alloy. Alloy is produced by casting on a cooled rotating roller.

EFFECT: high coercive force and good magnet characteristics of alloy.

7 cl, 25 dwg, 1 tbl, 1 ex

Description

FIELD OF TECHNOLOGY

[0001] The present invention relates to an R-T-B type alloy, a fine powder for an R-T-B type rare earth permanent magnet, and an R-T-B type rare earth permanent magnet. In particular, the present invention relates to an R-T-B type alloy and a fine powder for an R-T-B type rare earth permanent magnet, which can provide an R-T-B type rare earth permanent magnet with excellent coercive force.

BACKGROUND

[0002] R-T-B type magnets have been used for hard disk drives (HDD), magnetic resonance imaging (MRI), various types of engines, and the like. due to its high performance. The recent increase in demand for energy saving in addition to increasing the heat resistance of R-T-B type magnets has caused an increase in their use in engines, including car engines. R-T-B type magnets contain Nd, Fe and B as the main components, and therefore magnets of this type are collectively referred to as Nd-Fe-B type or R-T-B type magnets. In a magnet, R-T-B of type R is mainly Nd with a portion replaced by another rare earth element such as Pr, Dy and Tb; T is Fe with a part replaced by another transition metal such as Co and Ni; and B is boron and may be partially replaced by C or N.

[0003] An RTB type alloy that can be used in an RTB type magnet is an alloy in which the magnetic phase R ₂ T ₁₄ B, which promotes magnetization activity, is the main phase and coexists with a non-magnetic, rich in rare earth elements and low melting "rich in R" phase . Since this RTB type alloy is an active metal, it is usually melted or cast in a vacuum or in an inert gas. A sintered magnet is usually made from an cast RTB type alloy ingot according to the powder metallurgy process as follows. The alloy ingot is pulverized into an alloy powder with an average particle size of about 5 μm (d50 measured by a laser diffraction particle size distribution analyzer), molded by magnetic field compression, sintered at a high temperature of about 1000 to 1100 ° C in a sintering furnace, then subjected, if necessary, to heat treatment and machining, and then provided with a galvanic coating to increase corrosion resistance, thereby completing the production of sintered magnet.

[0004] In the sintered R-T-B type magnet, the rich R phase plays the following important roles:

1) forms a liquid phase during sintering due to the low melting point and thereby contributes to a high compaction of the magnet and, in turn, increase the magnetization;

2) eliminates irregularities on the grain boundary and thereby leads to a decrease in the nucleation centers of the inverse magnetic domains and an increase in the coercive force; and

3) magnetically isolates the main phase and thereby increases the coercive force.

Accordingly, if the R-rich phase in the molded magnet is in a poorly dispersed state, this causes local sintering defects or a decrease in magnetism. Therefore, it is important that the rich R phase is uniformly dispersed in the molded magnet. The distribution of the R-rich phase in the sintered R-T-B type magnet is significantly affected by the texture of the starting material of the R-T-B type alloy.

[0005] Another problem encountered when casting an R-T-B type alloy is the formation of α-Fe in the cast alloy. α-Fe has deformability and remains crushed in the grinder, and this not only reduces the grinding efficiency during grinding of the alloy, but also affects fluctuations in the composition or particle size distribution before and after grinding. If α-Fe still remains in the magnet after sintering, a decrease in the magnetic characteristics of the magnet occurs. Accordingly, in cases where necessary, the alloy was previously subjected to high temperature homogenization treatment for a long period of time in order to remove α-Fe. However, α-Fe is present as a peritectic nucleus and therefore its removal requires solid-phase diffusion over a long period of time. In the case of an ingot having a thickness of several centimeters and a rare earth content of 33% or less, removal of α-Fe is virtually impossible.

[0006] In order to solve the problem that α-Fe is formed in an RTB type alloy, a tape casting method (simply referred to as the "LR method") has been developed and used to cast an alloy ingot with an increased cooling rate. The LR method is a method of crystallizing an alloy by rapid cooling, in which the molten alloy is poured onto a copper roll, the inside of which is water-cooled, and a plate of about 0.1 to 1 mm is obtained. In the LR method, the molten alloy is supercooled to a temperature at which the main R ₂ T ₁₄ B phase is formed, or even lower, so that the R ₂ T ₁₄ B phase can be formed directly from the molten alloy, and α-Fe formation can be suppressed. In addition, in the LR method, a finely dispersed microstructure is created in the alloy, so that an alloy having a microstructure allowing fine dispersion of the R-rich phase can be obtained. The rich R phase expands upon reaction with hydrogen in a hydrogen atmosphere and becomes a brittle hydride. Using this property, fine crushing commensurate with the degree of dispersion of the rich R phase can be proposed. When the alloy is pulverized by this hydrogenation step, a large number of fine cracks are formed due to the initiation of hydrogenation to crack the alloy and, therefore, very good milling is achieved. The internal R-rich phase in the alloy obtained by the LR method is thus finely dispersed, and this leads to a good dispersion of the R-rich phase also in the magnet after grinding and sintering, thereby increasing the magnetic characteristics of the magnet (see, for example, patent document 1 )

[0007] Alloy plates obtained by the LR method are also excellent in terms of uniformity of microstructure. The uniformity of the microstructure can be compared by the diameter of the crystalline grain or the dispersed state of the R-rich phase. In the case of the alloy plate obtained by the LR method, a quenched crystal is sometimes formed on the side of the alloy plate (hereinafter referred to as the "surface side of the mold") facing the casting roll, but in general a properly fine homogeneous texture can be obtained, which is formed by crystallization through rapid cooling. As described above, in the R-T-B type alloy obtained by the LR method, the R-rich phase is finely dispersed and α-Fe formation is also suppressed, so the R-T-B type alloy has an excellent microstructure for making a sintered magnet.

[0008] The distribution of Dy, contributing to an increase in coercive force, significantly affects the characteristics of the magnet, especially the relationship between the coercive force and the distribution of elements in the microstructure of the magnet. For example, it was previously reported that the coercive force should be high when Dy is distributed close to the grain boundary phase (see, for example, Patent Document 2).

More specifically, it was previously reported that the coercive force should be high when Dy is present in the main phase (see, for example, Patent Document 3 and Non-Patent Document 1).

Additionally, since there is a definite relationship between the characteristics of the magnet and alloy manufacturing methods, alloy manufacturing methods have also been improved along with improvements in magnet characteristics. For example, a method for controlling microstructures is known (see, for example, Patent Document 4) and a method for controlling microstructures by treating the surface of a casting roll to a predetermined roughness (see, for example, Patent Documents 5 and 6).

Patent Document 1: Japanese Unexamined Patent Application, First Publication No. Hei 5-222488.

Patent Document 2: Unexamined Japanese Patent Application, First Publication No. Hei 5-21219.

Patent Document 3: WO 2003/001541.

Patent Document 4: WO 2005/031023.

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

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

Non-Patent Document 1: Hiroyuki TOMIZAWA, Journal of the Japan Society of Powder and Powder Metallurgy, March, 2005, vol. 52. issue 3, pp. 158-163.

SUMMARY OF THE INVENTION

Problems Solved by the Invention

[0009] However, in recent years, R-T-B type rare earth permanent magnets with even higher performance have been required, and thus, there is an increasing need for further improvements in magnetic characteristics, such as coercive force, R-T-B type rare earth permanent magnets.

The present invention was created in view of the foregoing circumstances, and it is an object of the present invention to provide an R-T-B type alloy as a starting material for a rare earth permanent magnet having excellent magnetic characteristics.

Another object of the present invention is to provide a fine powder for an R-T-B type rare earth permanent magnet made from the aforementioned R-T-B type alloy and an R-T-B type rare earth permanent magnet.

Means to Solve These Problems

[0010] The present inventors have carried out a detailed inspection of the texture of a Dy-containing RTB type alloy to be used to make RTB type rare earth permanent magnets in order to investigate the relationship between the texture state and magnetic characteristics. The authors of the present invention has also confirmed the fact that if the Dy-containing RTB type alloy includes a Dy-rich region enriched with Dy, in addition to the main phase formed from the R ₂ T ₁₄ B phase, and the R-rich phase enriched in R, then the RTB type rare earth permanent magnet obtained by molding / sintering a fine powder, which is made from plates of this RTB type alloy, will have excellent magnetic characteristics, such as coercive force. The present invention has been made based on these discovered facts.

[0011] That is, the present invention provides the following.

(1) An RTB type alloy (where R is at least one element selected from Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb and Lu; T is a transition metal that contains 80% by mass or more of Fe; and B is containing 50% by mass or more of boron (B) and also contains at least one element of C and N within the range from 0 to less than 50% by mass), which is the starting material for use in a rare earth permanent magnet and contains at least Dy containing a main phase, such as an R ₂ T ₁₄ B phase, having a magnetic properties, a rich R phase, which is relatively enriched in R compared to its share in the total alloy composition, and a rich Dy region, which is formed near the rich R phase and relatively enriched in Dy compared to the aforementioned share in the composition.

(2) An R-T-B type alloy as described in paragraph (1), wherein the Dy concentration is lower in the main phase than in the Dy rich region and lower in the R rich than in the main phase.

(3) An R-T-B type alloy as described in (1) or (2), the alloy being a plate having an average thickness of 0.1 to 1 mm made by a tape casting process.

[0012] (4) A method for manufacturing an RTB alloy of the type described in any of the above (1) to (3), comprising the steps of: making a plate having an average thickness of 0.1 to 1 mm, and supplying the molten alloy on a cooling roll with an average speed of 10 g / s or more per 1 cm of width.

(5) A method of manufacturing an R-T-B type alloy described in paragraph (4), characterized in that the R-T-B type alloy plate rolled from the cooling roll is held at a temperature of from 600 to 900 ° C for 30 seconds or more.

[0013] (6) Fine powder for an RTB type rare earth permanent magnet that is made of an RTB type alloy described in any of the above (1) to (3), or an RTB type alloy made by a method for manufacturing an RTB type alloy described in paragraph (4) or (5).

(7) R-T-B type rare earth permanent magnet made from fine powder of the R-T-B type rare earth permanent magnet described in (6).

The results of the invention

[0014] The R-T-B type alloy of the present invention is formed close to the R-rich phase and has a rich Dy region relatively rich in Dy compared to its share in the total composition. Accordingly, a rare earth permanent magnet having a high coercive force and excellent magnetic characteristics can be obtained.

Also, the fine powder for the RTB type rare earth permanent magnet and the RTB type rare earth permanent magnet of the present invention are made from either an RTB type alloy of the present invention or an RTB type alloy made by a method of manufacturing an RTB type alloy of the present invention, and thus will have high coercive force and excellent magnetic characteristics.

BRIEF DESCRIPTION OF THE DRAWINGS

[0015] FIG. 1 is a photograph showing one example of an R-T-B type alloy of the present invention. The photograph was taken while observing in a scanning electron microscope (SEM) the cross section of an R-T-B type alloy plate.

Figure 2 is an electronic image of an alloy of R-T-B type shown in Figure 1.

Figure 3 is an x-ray image of Fe in the region corresponding to that shown in Figure 2.

Figure 4 is an X-ray image of Nd in the region corresponding to that shown in Figure 2.

FIG. 5 is an X-ray image of Dy in the region corresponding to that shown in FIG. 2.

6 is an X-ray image of Ga in the region corresponding to that shown in FIG. 2.

Fig.7 is an electronic image of an alloy of R-T-B type shown in Fig.1.

Fig. 8 is an X-ray image of Dy in the region corresponding to that shown in Fig. 7.

Fig.9 is an x-ray image of Fe in the region corresponding to that shown in Fig.7.

Figure 10 is an x-ray image of Nd in the region corresponding to that shown in Figure 7.

11 is a schematic front view showing the construction of an alloy manufacturing apparatus according to an embodiment of the present invention.

12 is a schematic front view showing a casting device provided in an alloy manufacturing apparatus.

13 is a schematic front view showing a heating device provided in an alloy manufacturing apparatus.

Fig. 14 is a schematic side view showing this heating device provided in the alloy manufacturing apparatus.

Fig. 15 is a schematic plan view showing opening - closing platforms and a storage vessel (container) that are provided in the alloy manufacturing apparatus.

Fig is a schematic front view illustrating the operation of the installation for the manufacture of alloy.

17 is a schematic front view illustrating the operation of an alloy manufacturing apparatus.

Fig. 18 is a schematic front view illustrating the operation of an alloy manufacturing apparatus.

Fig. 19 is a schematic front view illustrating the operation of an alloy manufacturing apparatus.

FIG. 20 is a schematic side view illustrating an operation of an alloy manufacturing apparatus. FIG.

21 is an electronic image of an R-T-B type alloy having no Dy rich regions.

Fig - x-ray image Dy in the area corresponding to that shown in Fig.21.

Fig - x-ray image of Fe in the region corresponding to that shown in Fig.21.

Fig.24 is an X-ray image of Nd in the region corresponding to that shown in Fig.21.

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

DESCRIPTION OF REFERENCE POSITIONS

[0016] 1 - Technological installation (installation for the manufacture of alloy); 2 - casting device; 3 - heating device; 4 - storage vessel; 4a - cooling plate; 5 - container; 6 - camera; 7 - funnel; 7a - release of the funnel; 21 - crushing device; 31 - heater; 31c is the open part; 33 - opening-closing platform; 33a - platform plate; 33b is an open-close system; 51 - belt conveyor (mobile device); L is the molten alloy; N is the plate of the cast alloy.

BEST MODE FOR CARRYING OUT THE INVENTION

[0017] Fig. 1 is a photograph of one example of an R-T-B-type alloy of the present invention. The photograph was taken while observing in a scanning electron microscope (SEM) the cross section of an R-T-B-type alloy plate. Note that in FIG. 1, the left side is the surface side of the mold. The R-T-B-type alloy shown in FIG. 1 is made by the LR method. This R-T-B-type alloy has a composition, in units of mass fractions, of 23% Nd, 9% Dy, 1% B, 1% Co and 0.2% Ga, and the rest is Fe. Please note that the composition of the alloy is RTB-type (in which R is at least one element selected from Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb and Lu; T is a transition metal that contains 80% by mass or more of Fe; and B is containing 50% by mass or more of boron and also contains at least one element of C and N within the range from 0 to less than 50 % by weight) of the present invention is not limited to the above specific composition, and the alloy can have any composition, provided that it is an RTB-type alloy containing m nshey least Dy.

[0018] The RTB type alloy shown in FIG. 1 consists of an R ₂ T ₁₄ B phase (main phase) and an R rich phase. 1, an R-rich phase is indicated by white, and an R ₂ T ₁₄ B phase (main phase) is indicated by gray. The R ₂ T ₁₄ B phase is mainly formed from a columnar crystal and partially formed from an equiaxed crystal. The average diameter of the crystalline grain of the phase R ₂ T ₁₄ B in the direction of the short axis is from 10 to 50 microns. In the R ₂ T ₁₄ B phase, at the grain boundary and inside the grain, there is a linear R rich phase extending along the direction of the long axis of the columnar crystal, or divided into separate particles or partially destroyed R rich phase. The rich R phase is a nonmagnetic phase having a low melting point and enriched with R compared to its share in the total composition. The average distance between R-rich phases is from 3 to 10 microns.

[0019] FIGS. 2-6 show the results of an element distribution analysis (digital display) of the R-T-B type alloy shown in FIG. 1 by electron probe microanalysis (EPMA) using a wavelength dispersive X-ray spectrometer (WDS).

Figure 2 shows an electronic image of an alloy of the RTB type shown in Figure 1. The R-rich phase is indicated in white, and the R ₂ T ₁₄ B phase (main phase) is indicated in gray.

Figure 3 shows an x-ray image of Fe in the region corresponding to that shown in Figure 2. From Fig.2 and 3 it is clear that the rich R phase contains less Fe compared with the main phase.

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

Figure 5 shows the x-ray image of Dy in the area corresponding to that shown in Figure 2. From Fig.2 and 5 it is clear that the rich R phase contains less Dy compared with the main phase.

Figure 6 shows an x-ray image of Ga in the region corresponding to that shown in Figure 2. From Fig.2 and 6 it is clear that the rich R phase contains more Ga compared with the main phase.

[0020] Figure 7-10 shows the results of the analysis of the distribution of elements (digital display) using an electron probe microanalyzer with field emission (FE-EPMA).

In Fig.7 shows an electronic image of the alloy RTB type shown in Fig.1. The rich R phase is indicated in white, and the R ₂ T ₁₄ B phase (main phase) is indicated in gray.

On Fig shows an x-ray image of Dy in the area corresponding to that shown in Fig.7. From Figs. 7 and 8, it is clear that the Dy rich region, which is relatively rich in Dy compared to the R rich phase and the main phase, is formed close to the R rich phase. In addition, it is clear from Fig. 8 that the concentration of Dy is lower in the main phase than in the rich Dy region, and even lower in the rich R phase.

In Fig.9 shows an x-ray image of Fe in the region corresponding to that shown in Fig.7. 7 and 9 it is clear that the rich R phase contains less Fe compared with the main phase.

Figure 10 shows an X-ray image of Nd in the region corresponding to that shown in Figure 7. 7 and 10 it is clear that the rich R phase contains more Nd compared with the main phase.

[0021] A manufacturing method

The R-T-B-type alloy of the present invention shown in FIG. 1 can be cast, for example, by the LR method, which uses the alloy manufacturing apparatus shown in FIG. 11.

Installation for the manufacture of alloy

11 is a schematic front view of a general configuration of an alloy manufacturing apparatus of the present embodiment.

The alloy fabrication apparatus 1 shown in FIG. 11 (hereinafter described as “process unit 1”) is typically equipped with a casting device 2, a crushing device 21 and a heating device 3. The heating device 3 includes a heater 31 and a container 5. A container 5 includes a storage vessel 4 and a group 32 of opening-closing platforms provided above the storage vessel 4.

The process unit 1 shown in FIG. 11 is equipped with a chamber 6. The chamber 6 includes a casting chamber 6a and a temperature-resistant storage chamber 6b, which is provided below the casting chamber 6a and which is 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 a temperature-resistant storage chamber 6b. In addition, a shutter 6e is provided in the temperature-resistant storage chamber 6b, and the temperature-resistant storage chamber 6b is closed by the shutter 6e, except when the container 5 is transported outside the temperature-resistant storage chamber 6b.

In addition, the casting device 2 is also equipped with a crushing device 21, and a funnel 7 is provided between the casting device 2 and the group 32 of open-close pads. The funnel 7 directs the cast alloy plate to the group 32 of open-close pads.

[0022] Foundry device

On Fig shows a schematic front view of the casting device 2, which is provided in the technological installation 1.

The casting device 2 shown in FIG. 12 includes a cooling roll 22 that casts molten alloy L into cast alloy M by rapidly cooling the molten alloy using a water cooling system (not shown); an intermediate casting device 23, which feeds the molten alloy L to the cooling roll 22, and a crushing device 21, which crushes the cast alloy M into plates N of the cast alloy. As shown in FIG. 12, the crushing device 21, for example, includes a pair of crushing rolls 21a.

[0023] Heating device

On Fig shows a schematic front view of the heating device 3, which is provided in the technological installation 1, on Fig shows its schematic side view, and Fig shows its schematic top view.

As shown in FIGS. 13-15, the heater 31 included in the heating device 3 has a heater cover 31a and a heater body 31b attached under the heater cover 31a. A heater cover 31a is provided therein so as to release heat released from the main body 31b towards the container 5 and to prevent heat from being released from the main body 31b into the casting chamber 6a. Also, if a lid 31a is provided in the heater, it can prevent damage to the main body 31b in the event of a sudden drop of a portion of molten alloy or cast alloy on it.

The heater 31 has an open portion 31c, and a funnel 7 outlet 7a is located in this open portion 31c. As a result, the cast alloy plate N, which passes through the funnel 7 and then falls from the casting device 2, can be supplied to a group 32 of opening-closing pads in the container 5, which is provided below the heater 31.

In addition, the heater 31, as shown in FIGS. 11 and 13, is located along the longitudinal direction of the conveyor belt 51 (the direction of movement of the container 5), which is provided inside the temperature-resistant storage chamber 6b.

This configuration makes it possible to uniformly maintain the temperature of the plate N of the cast alloy poured onto a group 32 of opening-closing pads in the container 5, even if the container 5 is moved inside the temperature-resistant storage chamber 6b.

[0024] A group 32 of opening-closing pads included in the heating device 3 is combined with a collecting vessel 4 to form a container 5. That is, the container 5 shown in FIGS. 13-15 is formed by a collecting vessel 4 and a group of 32 opening-closing platforms, which is provided above the storage vessel 4.

The opening-closing pad group 32 is equipped with a plurality of opening-closing pad 33, which are arranged along the direction of movement of the container 5. In addition, guide elements 52 are provided around the group 32 of opening-closing pad, and these guide elements 52 prevent the cast plates N from being scattered which fall through a funnel 7 into a temperature-resistant storage chamber 6b.

[0025] Each opening-closing platform 33 leaves a cast alloy plate N, which is supplied from the casting device 2, poured onto it in order to withstand the temperature with the heater 31 for a predetermined period of time, and dumps the cast alloy plate N into the storage vessel 4 after the exposure time. Each opening-closing platform 33 is equipped with a platform plate 33a and an opening-closing system 33b that opens or closes the platform plate 33a. Each open-close system 33b has a rotary axis 33b ₁ attached to one side of the pad plate 33a, and a drive unit (not shown) that rotates the rotary axis 33b ₁ . Each drive unit can freely rotate the rotary axis 33b ₁ , so that the angle of inclination of each platform plate 33a can be adjusted separately. The tilt angle of each pad of the pad 33a can be set arbitrarily in a range from 0 ° (when the pad of the pad 33a is horizontal (the position shown in FIG. 13 by a dotted dotted line)) to about 90 ° in a clockwise direction (when the pad of the pad 33a is almost vertical (position shown in Fig. 13 by a solid line)).

[0026] Thus, the opening-closing platform 33 can leave the cast alloy plate N poured onto the platform plate 33a for a predetermined temperature holding time by activating the opening-closing system 33b, and then can drop the cast alloy plate N down into the storage vessel 4, increasing the angle of the plate 33a of the site.

In addition, the opening-closing platform 33 can act as a cover for the storage vessel 4 and this prevents heat from reaching the heater 31 of the storage vessel 4, thereby preventing heating of the inside of the storage vessel 4. Furthermore, a plurality of cooling plates are provided inside the storage vessel 4 4a.

[0027] In addition, as shown in FIGS. 13 and 14, the container 5 is mounted on the conveyor belt 51 (mobile device). The conveyor belt 51 allows the container 5 to move to the left or right side in FIG. 13.

[0028] Casting the alloy

All Figs. 16-19 are schematic front views illustrating the operation of the alloy manufacturing apparatus.

As shown in FIG. 16, the container is first moved to where an opening-closing platform 33A (located on the left edge of the group of 32 opening-closing platforms 32 in this figure) is located immediately below the outlet 7a of the funnel 7. Also, all opening-closing platforms 33 are set to a closed state.

Then prepare plates N of the cast alloy by actuating the casting device 2 shown in Fig.12. The molten alloy L is first prepared in a melting device (not shown). The temperature of the molten alloy L varies depending on the types of alloy composition, but it is set within the range from 1300 ° C to 1500 ° C. Prepared molten alloy L is transported to a casting device 2, while it is contained in a refractory crucible 24.

Then, the molten alloy L is supplied from the refractory crucible 24 to the intermediate casting device 23 and then fed from the intermediate casting device 23 to the cooling roll 22, whereby the molten alloy L is solidified (crystallized) to obtain a cast alloy M. After this, the cast alloy M is displaced from the cooling the roll 22 is on the opposite side from the intermediate casting device 23 and is inserted between two rotating crushing rolls 21a, so that the cast alloy M is crushed into the plate N of the cast alloy.

[0029] The average feed rate of the molten alloy to the cooling roll 22 is 10 g / s or more, preferably 20 g / s or more, more preferably 25 g / s or more, per 1 cm of width, and even more preferably 100 g / s or less per 1 cm of width. If the feed rate of the molten alloy L is less than 10 g / s, then the molten alloy L may not spread by the wetting thin layer and not be distributed on the cooling roll 22, but instead may compress due to the viscosity of the molten alloy L or the wettability of the surface of the cooling roll 22 and may a deviation in alloy quality appears. On the other hand, if the average feed rate of the molten alloy to the cooling roll 22 exceeds 100 g / s per 1 cm of width, then cooling on the cooling roll 22 may be insufficient and may cause microstructure coarsening, precipitation of α-Fe or the like.

[0030] The average cooling rate of the molten alloy on the cooling roll 22 is preferably from 100 to 2000 ° C / sec. An average cooling rate of 100 ° C / sec or more will be satisfactory to prevent the release of α-Fe or coarsening of the texture of the R-rich phase or the like. On the other hand, if the average cooling rate is 2000 ° C / sec or less, the degree of subcooling will not be excessive, and the cast alloy plate may be supplied to the heating device 3 at an appropriate temperature. In addition, the cast alloy plate does not cool too much and therefore does not require a reheating process. Note that the average cooling rate is determined by dividing the difference between the temperature immediately before the contact of the molten alloy with the cooling roll and the temperature when separated from the cooling roll by the time during which the molten alloy is in contact with the cooling roll.

[0031] The average temperature of the cast alloy M when separated from the cooling roll 22 varies slightly due to the subtle difference in the degree of contact between the cast alloy M and the cooling roll 22, fluctuations in the thickness of the cast alloy M, or the like. The average temperature of the cast alloy M when separated from the cooling roll can be obtained, for example, by scanning the surface of the alloy in the width direction using a radiation pyrometer from the beginning to the end of the casting, thus measuring the temperature and averaging the measured values.

[0032] The average temperature of the cast alloy M when separated from the cooling roll 22 is preferably 100-500 ° C lower, more preferably 100-400 ° C lower than the solidification temperature of the phase R ₂ T ₁₄ B in the equilibrium state of the molten alloy. It is generally accepted that the melting point of the R ₂ T ₁₄ B phase is 1150 ° C in the ternary Nd-Fe-B system, but varies in accordance with the substitution of Nd with other rare earth elements, the replacement of Fe with other transition elements, and the type and amount of any additive elements added. If the difference between the average temperature of the cast alloy M during separation from the cooling roll 22 and the solidification temperature of the phase R ₂ T ₁₄ B in the equilibrium state of the cast alloy M is less than 100 ° C, this may correspond to an insufficient cooling rate. On the other hand, if this difference exceeds 500 ° C, the supercooling of the molten alloy may become excessively large due to the too high cooling rate.

[0033] The average temperature of the cast alloy M when separated from the cooling roll 22 also varies within the same casting (smelting) stage, and if the width of such a spread is large, this can cause fluctuation in the microstructure or quality. Therefore, the width of the temperature variation within the melting is suitably less than 200 ° C, preferably 100 ° C or less, more preferably 50 ° C, even more preferably 20 ° C.

[0034] The plate N of the cast alloy preferably has an average thickness of 0.1 to 1 mm. If the average plate thickness is less than 0.1 mm, then the solidification rate may be excessively high and the R-rich phase may be dispersed too finely. On the other hand, if the average plate thickness exceeds 1 mm, the solidification rate may decrease and this may cause a decrease in the dispersion of the rich R phase, precipitation of α-Fe or the like.

[0035] Then, as shown in FIG. 16, the plates N of the cast alloy are delivered to the heating device 3 by passing through a funnel 7 and stacked (poured) onto that opening-closing platform 33, which is located directly under the outlet 7a of the funnel 7. During At this time, the heater 31 is turned on and the plates N of the cast alloy are held at a predetermined temperature or are heated by the heater 31 immediately after they are laid on the opening-closing platform 33A.

The number of plates N of the cast alloy laid on the open-close platform 33A can be adjusted accordingly according to the area of the platform plate 33a. However, since the plates N of the cast alloy are continuously fed from the casting device 2, over time they will drain from the opening-closing platform 33, although this also depends on the feed rate. For this reason, the container 5 moves to the left side in the drawing, as shown in FIG. 17, when the stacked number of plates N of the cast alloy reaches a predetermined value with respect to the opening-closing platform 33A. Then, another opening-closing pad 33B adjacent to the opening-closing pad 33A on the right side is placed directly below the outlet 7a of the funnel 7, followed by laying the plates N of the cast alloy onto the opening-closing pad 33B. After that, in the same way, the container 5 is moved in accordance with the preparation of the plates N of the cast alloy and these plates N of the cast alloy are stacked sequentially on the opening-closing platforms 33C-33E.

[0036] The cast alloy plates N, laid on each of the opening-closing platforms 33A-33E, are held at a predetermined temperature or heated by a heater 31. It is preferable that the holding temperature be lower than the temperature of the plate N when separated from the cooling roll (separation temperature ), and specifically, it is more preferable to be within a range from (compartment temperature - 100 ° C) to a compartment temperature, and more preferably within a range from (compartment temperature - 50 ° C) to the temperature of the compartment. More specifically, the holding temperature is preferably within the range of 600 ° C to 900 ° C. When the holding temperature is 600 ° C or more, the coercive force of the R-T-B-type alloy can be sufficiently increased. Also, when the holding temperature is 900 ° C or less, precipitation of α-Fe can be prevented and the microstructure, such as the R-rich phase, can be prevented from being coarse.

In addition, when the separation temperature decreases for any reason, the plates N of the cast alloy can be heated and held at a predetermined temperature by setting the holding temperature above the separation temperature. Preferably, the heating range is within 100 ° C, and more preferably within 50 ° C. If the heating range is too large, manufacturing efficiency will decrease. It should be noted that the coercive force can be increased even if the plates are held at 1000 ° C. However, this temperature makes the microstructure coarse. In addition, the distribution of particles or the fluidity of the fine powder during fine crushing and sintering temperature can adversely change. Accordingly, when they are held at a temperature of 1000 ° C, it is necessary to consider its effect on subsequent processes.

[0037] Furthermore, the temperature holding time is preferably 30 seconds or more, more preferably from 30 seconds to about several hours, and most preferably from 30 seconds to about 30 minutes. If the temperature holding time is 30 seconds or more, then the coercive force can be sufficiently increased. That is, the cast alloy plates may be heat treated by holding the temperature for several hours, but preferably the holding time is 30 minutes or less in terms of manufacturing efficiency.

[0038] Then, as shown in FIG. 18, the container 5 is then moved relative to the rest of the opening-closing pads 33F-33J in accordance with the preparation of the N alloy plates in the same manner, whereby the N alloy plates are successively stacked on each from the opening-closing platforms 33F-33J. As for the plates N of the cast alloy laid on the opening-closing platforms 33A-33D, they are successively discharged into the storage vessel 4 by sequentially putting each opening-closing platform in the open state, as shown in Fig. 18, when the predetermined holding time passes temperature. As soon as the plates N of the cast alloy are discharged into the storage vessel 4, the heat of the heater 31 no longer reaches the plates N of the cast alloy, so that the heat treatment by temperature exposure is completed.

As described above with reference to FIG. 17, the plates N of the cast alloy are successively poured onto each opening-closing platform, and as a result, various opening-closing platforms have time differences at the starting point of the heat treatment starting by holding the temperature relative to the plates N of the cast alloy on the opening- lockable areas. Therefore, it is preferable that the plates N of the cast alloy are sequentially discharged into the storage vessel 4 by sequentially switching each opening-closing pad to the open state in order to fix the temperature holding time with respect to the plates N of the cast alloy at each opening-closing pad.

The plate N of the cast alloy that is discharged into the storage vessel 4 is in contact with the cooling plate 4a, whereby heat is absorbed into the cooling plate 4a and the plate N of the cast alloy is thereby cooled.

[0039] Figs. 19 and 20 show a state in which all opening-closing platforms 33 are in the open state and the plates N of the cast alloy are stored in the storage vessel 4. If, after this, the casting and crushing processes are then carried out by the casting device 2 again, then the container 5 can be moved to the right in these figures, while all the opening-closing pads 33 are set to the closed state and the plates N of the cast alloy are successively poured onto each opening-closing pad 33 in accordance with and with the preparation of the flakes N of the cast alloy. On the contrary, if the casting and crushing processes by the casting device 2 are completed, then all opening-closing platforms 33 are switched to the closed state to prevent the heat of the heater 31 from reaching the storage vessel 4. Then, the shutter 6e of the temperature-resistant storage chamber 6b is opened and the container 5 is transported from the camera 6 outward to assemble the plates N of the cast alloy, thereby completing the manufacture of the plates N of the cast alloy.

[0040] Cooling rate

Next, the cooling rate in the manufacture of an R-T-B-type alloy will be described.

In the present invention, the cooling rate was adjusted so that the following cooling rates were achieved from the solidification point of the main phase (about 1170 ° C), which is the temperature immediately after solidification, to 600 ° C, which is below the solidification point of the R-rich phase (about 700 ° C) .

That is, the cooling rate of the R-T-B-type alloy from 1000 ° C to 850 ° C is set within the range from 100 to 300 ° C / sec. If the cooling rate from 1000 ° C to 850 ° C is higher than the above range, then Dy may not sufficiently diffuse into the main phase. On the other hand, if the cooling rate is lower than the above range, then Dy may excessively diffuse, making it impossible to form a rich Dy region in the main phase.

It is also preferred that the cooling rate of the R-T-B-type alloy from the solidification point of the main phase to 1000 ° C is set within the range of 300 to 2000 ° C / sec. By setting the cooling rate from the solidification point of the main phase to 1000 ° C within the above range, an R-T-B-type alloy with a rich Dy region is obtained with high productivity.

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

[0041] The R-T-B-type alloy and the R-T-B-type alloy plates of the present embodiment are formed close to the R-rich phase and have a rich Dy region that is relatively rich in Dy compared to its share in the total composition. Accordingly, a rare earth permanent magnet having a high coercive force and excellent magnetic characteristics can be obtained from them.

In other words, the R-T-B-type alloy of the present embodiment has a higher coercive force than, for example, the R-T-B-type alloy shown in FIGS. 21-24 without Dy-rich regions.

[0042] FIGS. 21-24 show, as one example, the results of an element distribution analysis (digital display) of an R-T-B type alloy without Dy-rich regions using an electron probe electron-probe microanalyzer (FE-EPMA). The R-T-B-type alloy shown in FIGS. 21-24 is made by the LR method. This R-T-B-type alloy has a composition, in units of mass fractions, of 23% Nd, 9% Dy, 1% B, 1% Co and 0.2% Ga, and the rest is Fe.

[0043] FIG. 21 is an electronic image of an RTB type alloy having no Dy rich regions. The rich R phase is indicated in white, and the R ₂ T ₁₄ B phase (main phase) is indicated in gray.

On Fig shows an x-ray image of Dy in the area corresponding to that shown in Fig.21. It is clear from Figs. 21 and 22 that this R-T-B-type alloy does not have any Dy rich region that would be enriched with Dy compared to the main phase, and that the concentration of Dy is lower in the R rich than in the main phase.

On Fig shows an x-ray image of Fe in the region corresponding to that shown in Fig.21. From Fig.21 and 23 it is clear that the rich R phase contains less Fe compared with the main phase.

On Fig shows an x-ray image of Nd in the region corresponding to that shown in Fig.21. From Fig.21 and 24 it is clear that the rich R phase contains more Nd compared with the main phase.

[0044] The manufacture of the R-T-B type rare earth permanent magnet.

To produce the R-T-B-type rare earth permanent magnet of the present invention, a fine powder for R-T-B-type rare earth permanent magnets is first made from the R-T-B-type alloy of the present invention. This fine powder for the RTB type rare earth permanent magnets of the present invention is obtained, for example, by performing hydrogen cracking of a plate formed from an RTB type alloy of the present invention by absorbing hydrogen and then converting the plate into powder using a grinder such as a jet mill . In this case of hydrogen cracking, for example, the step of absorbing hydrogen while holding the plate in a hydrogen atmosphere under a predetermined pressure is preferably performed in advance.

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

[0045] The fine powder for the R-T-B type rare earth permanent magnet and the R-T-B type rare earth permanent magnet according to the present embodiment are made from the R-T-B type alloy of the present invention. Accordingly, they will have a high coercive force and excellent magnetic characteristics.

Example 1

[0046] The starting metals constituted to obtain the alloy composition, in units of mass fractions, 23% Nd, 9% Dy, 0.98% B, 1% Co and 0.2% Ga, and the rest is Fe, weighed and then melted in an alumina crucible in an argon gas atmosphere at 1 atm. using a high frequency melting furnace to produce a molten alloy. Then this molten alloy was fed into the casting device in the processing unit shown in FIG. 11 and cast by the LR method. The rotation speed of the cooling roll during casting was 1.3 m / s, the average feed rate of the molten alloy to the cooling roll was 30 g / s per 1 cm of width, and the average temperature of the cast alloy ingot during separation from the cooling roll was 850 ° C.

The cooling rate of this alloy was 700 ° C / sec from the solidification point of the main phase to 1000 ° C, 200 ° C / sec from 1000 ° C to 850 ° C, and 50 ° C / sec from 850 ° C to 780 ° C. After that, the alloy was kept at a temperature of about 780 ° C for 300 seconds on an open-close platform using the processing unit shown in Fig. 11, and then cooled to 600 ° C or less with a cooling rate of 0.1 ° C / s, having obtained RTB-type alloy plates according to Example 1. The average alloy thickness at this stage was 0.3 mm.

Example 2

[0047] The molten alloy was made using the same starting metals and the same setup as in Example 1. Then, the obtained molten alloy was cast using the same casting device as in Example 1. The rotation speed of the cooling roll during casting was 0, 87 m / s, the average feed rate of the molten alloy to the cooling roll was 30 g / s per 1 cm of width, and the average temperature of the cast alloy ingot when separated from the cooling roll was 880 ° C.

The cooling rate of this alloy was 700 ° C / sec from the solidification point of the main phase to 1000 ° C, 200 ° C / sec from 1000 ° C to 850 ° C, and 10 ° C / sec from 850 ° C to 780 ° C. After that, the alloy was cooled to 600 ° C or less with a cooling rate of 0.1 ° C / s using the processing unit shown in Fig. 11 without using an opening-closing pad, having obtained RTB-type alloy plates according to Example 2. Average thickness the alloy at this stage was 0.45 mm.

[0048] The RTB-type alloy plates obtained in Examples 1 and 2 were subjected to element distribution analysis (digital display) (surface analysis) using an electron probe microanalyzer equipped with a wavelength dispersive X-ray spectrometer (WDS-EPMA), and electron probe microanalyzers with field emission (FE-EPMA). As a result, it was found that both R-T-B-type alloy plates obtained in Examples 1 and 2 formed a Dy-rich region that was enriched in Dy compared to the R-rich and main phase, near the R-rich phase. In addition, in both R-T-B-type alloy plates obtained in Examples 1 and 2, the Dy concentration was lower in the main phase than in the Dy rich region and even lower in the R rich.

Comparative Example 1

[0049] The molten alloy was made using the same starting metals and the same setup as in Example 1. Then, the obtained molten alloy was cast using the same casting device as in Example 1, making RTB-type alloy plates according to comparative example 1. The rotation speed of the cooling roll during casting was 0.65 m / s, the average feed rate of the molten alloy to the cooling roll was 15 g / s per 1 cm of width, and the average temperature of the cast alloy ingot when separated from the cooling of the roll was 700 ° C.

The cooling rate of this alloy was 700 ° C / sec from the solidification point of the main phase to 1000 ° C, 400 ° C / sec from 1000 ° C to 700 ° C, and 10 ° C / sec from 700 ° C to 600 ° C. After that, the alloy was cooled to 600 ° C or less with a cooling rate of 0.1 ° C / s using the processing unit shown in Fig. 11 without using an opening-closing platform. The average alloy thickness at this stage was 0.30 mm.

[0050] The R-T-B-type alloy plates obtained in comparative example 1 were subjected to element distribution analysis (digital display) (surface analysis) using WDS-EPMA and FE-EPMA. As a result, it was found that the R-T-B-type alloy plates obtained in comparative example 1 did not form any Dy-rich regions that would be enriched with Dy compared to its share in the total composition. One possible reason for this result could be that in comparative example 1, the temperature of the cast ingot during separation from the cooling roll was low, and the alloy cooled too quickly on the cooling roll, which made the cooling rate of the alloy from 1000 ° C to 700 ° C too high . Accordingly, Dy and Nd may not have diffused sufficiently, and their concentration gradients did not form.

[0051] Then, magnets were made as described below using the R-T-B-type alloy plates obtained in Examples 1 and 2 and Comparative Example 1.

The R-T-B-type alloy plates obtained in Examples 1 and 2 and Comparative Example 1 were first subjected to hydrogen cracking. Hydrogen cracking was carried out as follows. The plates of R-T-B-type alloys were forced to absorb hydrogen in a hydrogen atmosphere at a pressure of 2 atm, and then heated to 500 ° C in vacuum for dehydrogenation. Thereafter, 0.07% by weight of zinc stearate was added thereto, and the resulting mixture was pulverized using a jet mill using a nitrogen gas stream. The powder obtained by this conversion had an average grain size of about 5.0 μm as measured by laser diffraction.

[0052] Then, the obtained powder material was molded by compression molding in a transverse magnetic field at a molding pressure of 0.8 t / cm ² in an atmosphere of 100% nitrogen to obtain a powder compact. The resulting powder compact was heated in vacuo at 1.33 × 10 ^-5 hPa from room temperature and kept at 500 ° C for 1 hour and then at 800 ° C for 1 hour to remove zinc stearate and remaining hydrogen. Then, the resulting powder compact was heated to a sintering temperature of 1030 ° C and held there for 3 hours to obtain a sintered product. After that, the obtained sintered product was further heat treated for 1 hour at 800 ° C and then at 530 ° C in an argon atmosphere. As a result, in both Examples 1 and 2, 10 magnets were obtained, and in comparative example 1, 5 magnets were obtained.

The magnetic characteristics of the magnets obtained in examples 1 and 2 and comparative example 1 were measured by means of a direct current magnetization curve inductor BH. The results are shown in Table 1 and FIG. 25. On Fig shows a graph of the coercive force (Hcj) of the magnets manufactured in Examples 1 and 2 and comparative example 1. In this graph, the vertical axis indicates the coercive force for the respective examples, which are shown on the horizontal axis. In Fig.25, the symbol "○" indicates the coercive force of the magnets manufactured in examples 1 and 2, and the symbol "▲" indicates the coercive force of the magnets manufactured in comparative example 1.

[0053]

Table 1 Hcj (kE) Hk / Hcj (%) BH _max (MGSE) Br (kgf) Example 1 33.34 56.40% 32.76 11.54 34.19 54.87% 32.61 11.54 33.69 56.39% 33,20 11.62 33,99 56.73% 33.36 11.63 34.10 56.43% 33.27 11.63 33.35 55.58% 32.41 11.47 33.18 57.28% 33.27 11.61 33.44 57.95% 33.72 11.71 33.67 56.62% 33.08 11.58 33.13 57.53% 33.22 11.61 Example 2 32.61 58.71% 33.26 11.64 33.07 58.03% 32.96 11.58 33.31 58.49% 33.66 11.68 33.74 57.79% 33.15 11.61 33.01 58.33% 33.03 11.58 33.31 58.09% 32.82 11.60 33.31 58.84% 32.87 11.57 33.69 57.53% 33.00 11.60 33.47 59.14% 33.35 11.65 33.04 60.06% 33.22 11.66 Comparative Example 1 31.44 52.30% 32.22 11.45 31.73 52.22% 32.73 11.54 31.28 53.54% 32.73 11.53 31.38 53.08% 33,50 11.70 31,42 52.93% 32.92 11.57

[0054] In Table 1, “(BH) _max ” indicates the maximum energy product of the magnet, “Br” indicates the residual magnetic induction, “Hcj” indicates the coercive force, and “Hk / Hcj” indicates the squareness of the hysteresis loop.

[0055] As shown in Table 1 and Fig. 25, the magnets obtained in examples 1 and 2 had higher coercive forces "Hcj" compared to those for the magnets obtained in comparative example 1, which were made from alloy RTB- type without formed rich Dy areas. Such differences in the coercive force of the magnets were caused by the distribution in them of the concentration of elements, which arose when they were in the alloy state, and which continued to influence even after the alloys were crushed and sintered to make magnets. One possible reason for this is that since Dy-rich regions are present in the alloys of the present invention and they also remain in the crystalline grains of the magnets, only a small amount of Dy remains in the R-rich phase, which does not actually contribute to the improvement of the coercive force.

Claims (7)

1. RTB-type alloy, where R is at least one element selected from Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Tb, Ho, Er, Tm, Yb and Lu; T is a transition metal that contains 80% by mass or more of Fe and B is containing 50% by mass or more of boron (B) and also contains at least one element of C and N in the range from 0 to less than 50% by the mass, which is the starting material for use in a permanent magnet based on rare-earth elements and contains at least Dy, and said RTB-type alloy contains:
a main phase, such as a phase R ₂ T ₁₄ B having magnetic properties;
rich in R phase, which is enriched in R compared with its share in the total composition of the alloy; and
Dy-rich region, which is formed near the R-rich phase and enriched in Dy in comparison with its share in the composition. 2. The R-T-B-type alloy according to claim 1, wherein the concentration of Dy is lower in the main phase than in the rich Dy region and lower in the rich R phase than in the main phase. 3. The R-T-B-type alloy according to claim 1, wherein this alloy is a plate having an average thickness of 0.1 to 1 mm made by a tape casting method. 4. A method of manufacturing an RTB-type alloy according to claim 1, comprising the steps of:
make a plate having an average thickness of from 0.1 to 1 mm; and
the molten alloy is fed to the cooling roll at an average speed of 10 g / s or more per 1 cm of width. 5. The method according to claim 4, in which the R-T-B-type alloy plate rolled from the cooling roll is held at a temperature of from 600 to 900 ° C for 30 s or more. 6. Fine powder for an R-T-B-type rare earth permanent magnet, which is made of an R-T-B-type alloy according to any one of claims 1 to 3, or of an R-T-B-type alloy made by the method of manufacturing an R-T-B-type alloy according to claim 4 or 5. 7. R-T-B-type rare earth permanent magnet made of fine powder for the R-T-B-type rare earth permanent magnet according to claim 6.

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