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
An object of the present invention is to provide an R-T-B type alloy (wherein 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 one that contains 50% by mass or more of boron (B) and also contains at least one element of C and N within a range from 0 to less than 50% by mass) that contains at least Dy, as a raw material for a rare earth-based permanent magnet having excellent magnetic characteristics, and the R-T-B type alloy provided in the present invention includes a main phase such as an R2T14B phase for exhibiting magnetic properties; an R-rich phase that is relatively enriched with R compared to the overall alloy compositional ratio; and a Dy-rich region that is formed close to the R-rich phase and relatively enriched with Dy compared to the compositional ratio.
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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 that can provide an R-T-B type rare earth permanent magnet having excellent coercive force.
BACKGROUND ARTR-T-B type magnets have been used for hard disks (HD), magnetic resonance imaging (MRI), various types of motors and the like for their high-performance characteristics. A recent increase in demand for energy saving, in addition to enhancements in the heat resistance of R-T-B type magnets, has caused the usage rate in motors, including automobile motors, to increase.
R-T-B type magnets have Nd, Fe and B as the main components and thus, the magnets of this type are collectively called Nd—Fe—B type or R-T-B type magnets. In an R-T-B type magnet, R is primarily Nd with a part being replaced by another rare earth element such as Pr, Dy and Tb; T is Fe with a part being replaced by another transition metal such as Co and Ni; and B is boron and may be partially replaced by C or N. The R-T-B type alloy which can be used in an R-T-B type magnet is an alloy where a magnetic R2T14B phase contributing to the magnetization activity is the main phase and coexists with a non-magnetic, rare earth element-enriched and low-melting point R-rich phase. Since this R-T-B type alloy is an active metal, it is generally melted or cast in vacuum or in an inert gas. From the cast R-T-B type alloy ingot, a sintered magnet is usually produced by a powder metallurgy process as follows. The alloy ingot is ground into an alloy powder with an average particle size of about 5 μm (d50: measured by a laser-diffraction particle size distribution analyzer), press-shaped in a magnetic field, sintered at a high temperature of about 1,000 to 1,100° C. in a sintering furnace, then subjected to, if necessary, heat treatment and machining, and further plated for enhancing the corrosion resistance, thereby completing a sintered magnet.
In the R-T-B type sintered magnet, the R-rich phase plays the following important roles:
1) forming a liquid phase during the sintering by virtue of a low melting point and thereby contributing to high densification of the magnet and in turn, enhancement of the magnetization;
2) eliminating unevenness on the grain boundary and thereby yielding a reduction in the nucleation site of the reversed magnetic domain and an increase in the coercive force; and
3) magnetically isolating the main phase and thereby increasing the coercive force.
Accordingly, if the R-rich phase in the shaped magnet is in a poorly dispersed state, it incurs local failure of sintering or reduction of magnetism. Therefore, it is important that the R-rich phase is uniformly dispersed in the shaped magnet. The R-rich phase distribution in the R-T-B type sintered magnet is greatly affected by the texture of the raw material R-T-B type alloy.
Another problem encountered in casting an R-T-B type alloy is the production of α-Fe in the cast alloy. The α-Fe has deformability and remains in the grinder without being ground, and this not only decreases the grinding efficiency at the grinding of the alloy but also affects the compositional fluctuation or particle size distribution before and after grinding. If α-Fe still remains in the magnet after sintering, a reduction in the magnetic characteristics of the magnet results. Accordingly, an alloy has been heretofore subjected to a homogenization treatment at a high temperature for a long period of time, where necessary, to eliminate α-Fe. However, α-Fe is present as a peritectic nucleus, and therefore its elimination requires solid phase diffusion for 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, elimination of α-Fe is practically impossible.
In order to solve the problem that α-Fe is produced in the R-T-B type alloy, a strip casting method (simply referred to as an “SC method”) of casting an alloy ingot at a higher cooling rate has been developed and used. The SC method is a method of solidifying an alloy through rapid cooling, where a molten alloy is cast on a copper roll the inside of which is water-cooled, and a flake of about 0.1 to 1 mm is produced. In the method, the molten alloy is supercooled down to a temperature where the main R2T14B phase is produced or even lower, so that an R2T14B phase can be produced directly from a molten alloy and the formation of α-Fe can be suppressed. Furthermore, in the SC method, a fine microstructure is generated in the alloy, so that an alloy having a microstructure allowing for fine dispersion of an R-rich phase can be produced. The R-rich phase expands by reacting with hydrogen in a hydrogen atmosphere and becomes a brittle hydride. By utilizing this property, fine cracking commensurate with the dispersion degree of the R-rich phase can be introduced. When an alloy is pulverized through this hydrogenation step, a large number of fine cracks are produced by the hydrogenation trigger breakage of the alloy, and therefore very good grindability is attained. The internal R-rich phase in the alloy produced by the SC method is thus finely dispersed, and this leads to good dispersibility of the R-rich phase also in the magnet after grinding and sintering, thereby enhancing the magnetic characteristics of the magnet (see, for example, Patent Document 1).
The alloy flake produced by the SC method is also excellent in terms of microstructure homogeneity. The microstructure homogeneity can be compared by the crystal grain diameter or the dispersed state of the R-rich phase. In the case of an alloy flake produced by the SC method, a chill crystal is sometimes generated on the casting roll side of the alloy flake (hereinafter referred to as a “mold face side”), but an appropriately fine homogeneous texture yielded by the solidification through rapid cooling can be obtained as a whole. As described above, in the R-T-B type alloy produced by the SC method, the R-rich phase is finely dispersed and the formation of α-Fe is also suppressed, and thus the R-T-B type alloy has an excellent microstructure for the production of a sintered magnet.
The distribution of Dy contributing to enhancement of the coercive force greatly affects magnet characteristics, particularly the relationship between the coercive force and the element distribution in the microstructure of a magnet. For example, the coercive force has already been reported to be high when Dy is distributed close to the grain boundary phase (see, for example, Patent Document 2).
More specifically, the coercive force has also been reported to 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 magnet characteristics and the alloy production methods, the methods for producing alloys have also advanced along with the improvements of magnet characteristics. For example, a method for controlling microstructures (see, for example, Patent Document 4) and a method for controlling microstructures by processing the surface of a casting roll to a predetermined roughness (see, for example, Patent Documents 5 and 6) are known.
[Patent Document 1] Japanese Unexamined Patent Application, First Publication No. Hei 5-222488
[Patent Document 2] Japanese Unexamined Patent Application, First Publication No. 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] Hiroyuki TOMIZAWA, Journal of the Japan Society of Powder and Powder Metallurgy, March, 2005, vol. 52, issue 3, pp. 158-163.
DISCLOSURE OF THE INVENTION Problems to be Solved by the InventionHowever, in recent years, R-T-B type rare earth permanent magnets with even higher performance have been required and thus, demands for further improvements in the magnetic characteristics, such as the coercive force, of R-T-B type rare earth permanent magnets are increasing.
The present invention has been made in view of the above circumstances and an object of the present invention is to provide an R-T-B type alloy as a raw material for a rare earth-based 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 produced from the above R-T-B type alloy, and an R-T-B type rare earth permanent magnet.
Means for Solving the ProblemsThe present inventors have conducted a detailed observation of the texture of the R-T-B type alloy containing Dy to be used for producing R-T-B type rare earth permanent magnets to investigate the relationship between the texture state and the magnetic characteristics. Also, the present inventors have confirmed the fact that when the R-T-B type alloy containing Dy includes a Dy-rich region enriched with Dy in addition to the main phase formed from an R2T14B phase and an R-rich phase enriched with R, the R-T-B type rare earth permanent magnet obtained by shaping/sintering a fine powder, which is produced from the flakes of this R-T-B type alloy, will have excellent magnetic characteristics such as coercive force. The present invention has been accomplished based on these findings.
That is, the present invention provides the following.
(1) An R-T-B type alloy (in which R is at least one element selected from Sc, Y, La, Ce, Pr, Nd, Pm, Sm, Eu, Gd, Th, Ho, Er, Tm, Yb, and Lu; T is a transition metal that contains 80% by mass or more of Fe; and B is one that contains 50% by mass or more of boron (B) and also contains at least one element of C and N within a range from 0 to less than 50% by mass) which is a raw material for use in a rare earth-based permanent magnet and contains at least Dy, comprising a main phase such as an R2T14B phase for exhibiting magnetic properties, an R-rich phase that is relatively enriched with R compared to the overall alloy compositional ratio, and a Dy-rich region that is formed close to the R-rich phase and relatively enriched with Dy compared to the aforementioned compositional ratio.
(2) The R-T-B type alloy as described in (1), wherein Dy concentration is lower in the main phase than in the Dy-rich region and is lower in the R-rich phase than in the main phase.
(3) The R-T-B type alloy as described in (1) or (2), wherein the alloy is a flake having an average thickness from 0.1 to 1 mm produced by a strip casting method.
(4) A method for producing an R-T-B type alloy described in any one of the above (1) to (3), comprising the steps of: producing a flake having an average thickness from 0.1 to 1 mm, and supplying molten alloy to a cooling roll at an average rate of 10 g/sec or more per 1-cm width.
(5) The method for producing an R-T-B type alloy described in (4), characterized in that an R-T-B type alloy flake rolled out of the cooling roll is kept at a temperature from 600 to 900° C. for 30 seconds or more.
(6) A fine powder for an R-T-B type rare earth permanent magnet that is produced from the R-T-B type alloy described in any one of the above (1) to (3) or from the R-T-B type alloy produced by the method for producing an R-T-B type alloy described in (4) or (5).
(7) An R-T-B type rare earth permanent magnet produced from the fine powder for an R-T-B type rare earth permanent magnet described in (6).
EFFECTS OF THE INVENTIONThe R-T-B type alloy of the present invention is formed close to the R-rich phase and has a Dy-rich region relatively enriched with Dy compared to the overall compositional ratio. Accordingly, a rare earth permanent magnet having a high coercive force and excellent magnetic characteristics can be achieved.
Also, the fine powder for an R-T-B type rare earth permanent magnet and the R-T-B type rare earth permanent magnet of the present invention are produced from either the R-T-B type alloy of the present invention or the R-T-B type alloy produced by the method of the present invention for producing an R-T-B type alloy, and thus will have a high coercive force and excellent magnetic characteristics.
1: Production apparatus (apparatus for producing alloy); 2: Casting device; 3: Heating device; 4: Storage vessel; 4a: Cooling plate; 5: Container; 6: Chamber; 7: Hopper; 7a: Hopper outlet; 21: Crushing device; 31: Heater; 31c: Opening part; 33: opening-closing stage; 33a: Stage plate; 33b: Opening-closing system; 51: Belt conveyor (movable device); L: Molten alloy; N: flake of cast alloy
BEST MODE FOR CARRYING OUT THE INVENTIONThe R-T-B type alloy shown in
The R-T-B type alloy shown in
The R-T-B type alloy of the present invention shown in
An apparatus 1 for producing an alloy shown in
The production apparatus 1 shown in
Moreover, the casting device 2 is also equipped with the crushing device 21 and a hopper 7 is provided between the casting device 2 and the opening-closing stage group 32. The hopper 7 directs the cast alloy flake onto the opening-closing stage group 32.
[Casting Device]The casting device 2 shown in
As shown in
The heater 31 has an opening part 31c, and an outlet 7a of the hopper 7 is disposed in the opening part 31c. Consequently, the flake N of the cast alloy that passes through the hopper 7 and then falls down from the casting device 2 can be supplied to an opening-closing stage group 32 in the container 5 which is provided below the heater 31.
Moreover, the heater 31, as shown in
This configuration makes it possible to uniformly maintain the temperature of the flake N of the cast alloy mounted on the opening-closing stage group 32 in the container 5 even when the container 5 moves inside the temperature-maintaining storage chamber 6b.
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 with the storage vessel 4, and the opening-closing stage group 32 which is provided over the storage vessel 4.
The opening-closing stage group 32 is equipped with a plurality of opening-closing stages 33 that are disposed along the moving direction of the container 5. In addition, guide members 52 are provided around the opening-closing stage group 32, and the guide members 52 prevent the flake N of the cast alloy that drop through the hopper 7 from scattering into the temperature-maintaining storage chamber 6b.
Each opening-closing stage 33 leaves the flake N of the cast alloy, which are supplied from the casting device 2, mounted thereon to maintain the temperature with the heater 31 at a predetermined period, and drops the flake N of the cast alloy to the storage vessel 4 after the temperature holding time. Each opening-closing stage 33 is equipped with a stage plate 33a, and an opening-closing system 33b which opens or closes the stage plate 33a. Each opening-closing system 33b has a rotating shaft 33b, attached to one side of the stage plate 33a; and a driving unit (not shown in the figures), which rotates the rotating shaft 33b1. Each driving unit can freely rotate the rotating shaft 33b1 such that the inclination angle of each stage plate 33a can be controlled separately. The inclination angle of each stage plate 33a can be set anywhere in the range of 0° (where the stage plate 33a is horizontal (the position shown in
Thus, the opening-closing stage 33 can leave the flake N of the cast alloy mounted on the stage plate 33a during a predetermined temperature holding time by actuating the opening-closing system 33b, and then, can drop the flake N of the cast alloy down into the storage vessel 4 by making the inclination angle of the stage plate 33a larger.
In addition, the opening-closing stage 33 can function as a cover for the storage vessel 4 and this prevents the heat of the heater 31 from reaching the storage vessel 4, thereby preventing the inside of the storage vessel 4 from heating up. Also, a plurality of cooling plates 4a is provided inside the storage vessel 4.
Moreover, as shown in
All
As shown in
Then, flakes N of the cast alloy are prepared by actuating the casting device 2 shown in
The average molten alloy supply rate to the cooling roll 22 is 10 g/sec or more, preferably 20 g/sec or more, more preferably 25 g/sec or more, per 1-cm width, and even more preferably 100 g/sec or less per 1-cm width. If the supply rate of molten alloy L is less than 10 g/sec, the molten alloy L may not be thinly wetted and spread on the cooling roll 22 and instead may shrink because of the viscosity of the molten alloy L itself or wettability to the surface of the casting roll 22 and fluctuation of the alloy quality may be brought about. On the other hand, if the average molten alloy supply rate to the cooling roll 22 exceeds 100 g/sec per 1-cm width, cooling on the cooling roll 22 may be insufficient and may cause coarsening of the microstructure, precipitation of α-Fe, or the like.
The average cooling rate of the molten alloy on the cooling roll 22 is preferably from 100 to 2,000° C./sec. An average cooling rate of 100° C./sec or more will be satisfactory to prevent precipitation of α-Fe or texture coarsening of the R-rich phase or the like. On the other hand, if the average cooling rate is 2,000° C./sec or less, the degree of supercooling will not be excessive and the cast alloy flake can be supplied to the heating device 3 at an appropriate temperature. Moreover, the cast alloy flake does not cool too much, and thus requires no reheating process. Note that the average cooling rate is determined by dividing the difference between the temperature immediately before contact of the molten alloy with the cooling roll and the temperature on detaching from the cooling roll by the time for which the molten alloy is contacted with the cooling roll.
The average temperature of the cast alloy M on detaching from the cooling roll 22 subtly varies due to a fine difference in the degree of contact between the cast alloy M and the cooling roll 22, fluctuation of the thickness of the cast alloy M, or the like. The average temperature of the cast alloy M on detaching from the cooling roll can be obtained, for example, by scanning the alloy surface in the width direction with a radiation thermometer from the start to finish of casting, thereby measuring the temperature, and averaging the measured values.
The average temperature of the cast alloy M on detaching from the cooling roll 22 is preferably 100 to 500° C. lower, more preferably 100 to 400° C. lower, than the solidification temperature of the R2T14B phase in an equilibrium state of the molten alloy. The melting temperature of the R2T14B phase is acknowledged to be 1,150° C. in the Nd—Fe—B ternary system but varies according to the substitution of Nd by other rare earth elements, the substitution of Fe by other transition elements, and the kind and amount added of any additive elements. If the difference between the average temperature of the cast alloy M on detaching from the cooling roll 22 and the solidification temperature of the R2T14B phase in an 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 molten alloy may become excessively large due to too high a cooling rate.
The average temperature of the cast alloy M on detaching from the cooling roll 22 also varies within the same casting step (tap) and if the variation width is large, this may bring about fluctuation of the microstructure or quality. Therefore, the variation width of temperature within the tap is suitably smaller than 200° C., preferably 100° C. or less, more preferably 50° C., even more preferably 20° C.
The cast alloy flake N preferably has an average thickness of 0.1 to 1 mm. If the average thickness of the flake is less than 0.1 mm, the solidification rate may be excessively increased and the R-rich phase may be dispersed too finely. On the other hand, if the average thickness of the flake exceeds 1 mm, the solidification rate may decrease and this may causes a reduction in the dispersibility of the R-rich phase, precipitation of α-Fe, or the like.
Next, as shown in
The amount of flakes N of the cast alloy piled on the opening-closing stage 33A may appropriately be adjusted in accordance with the area of the stage plate 33a. However, since the flakes N of the cast alloy are continuously supplied from the casting device 2, they will overflow from the opening-closing stage 33A in time although it also depends on the supply rate. For this reason, the container 5 is moved to the left-hand side in the drawing as shown in
The flakes N of the cast alloy piled on each of the opening-closing stages 33A to 33E are kept at a predetermined temperature or heated with the heater 31. It is preferable that the holding temperature be lower than the temperature of the flake N when detaching from the cooling roll (detaching temperature), and specifically, it is preferably within a range of (the detaching temperature −100° C.) to the detaching temperature, and it is more preferably within a range of (the detaching temperature −50° C.) to the detaching temperature. More specifically, the holding temperature is preferably within a range of 600° C. to 900° C. When the holding temperature is 600° C. or more, the coercive force of an R-T-B type alloy can be sufficiently enhanced. Also, when the holding temperature is 900° C. or less, the deposition 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 detaching temperature declines for any reason, the flakes N of the cast alloy can be heated and kept at a predetermined temperature by setting the holding temperature higher than the detaching temperature. It is preferable that the heating range be within 100° C., and more preferably within 50° C. If the heating range is too large, the production efficiency will decline. It should be noted that the coercive force can be improved even when the flakes are kept at 1,000° C. However, such a temperature makes microstructure coarse. Furthermore, the particle distribution or the fluidity of the fine powder when they are finely crushed, and the sintering temperature may unfavorably change. Accordingly, when they are kept at 1,000° C., it is required to consider its influence to subsequent processes.
Furthermore, the temperature holding time is preferably 30 seconds or more, more preferably 30 seconds to about several hours, and most preferably 30 seconds to about 30 minutes. If the temperature holding time is 30 seconds or more, then the coercive force can be sufficiently enhanced. That is, the flakes of the cast alloy may be subjected to the temperature holding treatment for several hours, but the temperature holding time is preferably 30 minutes or less in terms of the production efficiency.
Next, as shown in
As described above by referring to
The flakes N of the cast alloy that dropped into the storage vessel 4 are in contact with the cooling plate 4a, whereby the heat is absorbed into the cooling plate 4a, and the flakes N of the cast alloy are consequently cooled down.
Next, the cooling rate on producing the R-T-B type alloy will be described.
In the present invention, the cooling rate was controlled so as to achieve the following cooling rates from the solidifying point of the main phase (around 1,170° C.), which is a temperature immediately after the solidification, to 600° C., which is lower than the solidifying point of the R-rich phase (around 700° C.).
That is, the cooling rate of R-T-B type alloy from 1,000° C. to 850° C. is set within a range of 100 to 300° C./sec. If the cooling rate from 1,000° C. to 850° C. is higher than the above range, Dy may not sufficiently diffuse into the main phase. On the other hand, if the cooling rate is lower than the above range, Dy may diffuse excessively, making it impossible to form a Dy-rich region in the main phase.
Also, it is preferable that the cooling rate of R-T-B type alloy from the solidifying point of the main phase to 1,000° C. be set within a range of 300 to 2,000° C./sec. By setting the cooling rate from the solidifying point of the main phase to 1,000° C. within the above range, an R-T-B type alloy having a Dy-rich region is obtained with high productivity.
In addition, it is preferable that the cooling rate of R-T-B type alloy from 850° C. to 600° C. be 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 diffuse into the adjacent main phase sufficiently. Accordingly, an R-T-B type alloy having a Dy-rich region and even higher coercive force can easily be produced.
The R-T-B type alloy and the flakes of the R-T-B type alloy of the present embodiment are formed close to the R-rich phase and have a Dy-rich region that is relatively enriched with Dy compared to the overall compositional ratio. Accordingly, a rare earth permanent magnet having a high coercive force and excellent magnetic characteristics can be achieved from them.
In other words, the R-T-B type alloy of the present embodiment has a higher coercive force compared to, for example, the R-T-B type alloy shown in
For producing 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 produced from the R-T-B type alloy of the present invention. The fine powder for R-T-B type rare earth permanent magnets of the present invention is obtained, for example, by a method of performing hydrogen cracking of a flake formed of the R-T-B type alloy of the present invention by hydrogen absorption and then pulverizing the flake by using a grinder such as a jet mill. In the hydrogen cracking here, for example, a hydrogen absorption step of keeping the flake 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 is, for example, press-shaped by a shaping 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.
The fine powder for an R-T-B type rare earth permanent magnet and the R-T-B type rare earth permanent magnet of the present embodiment are produced from the R-T-B type alloy of the present invention. Accordingly, they will have a high coercive force and excellent magnetic characteristics.
EXAMPLE 1Starting metals formulated to have an alloy composition of, in terms of mass ratio, 23% of Nd, 9% of Dy, 0.98% of B, 1% of Co, and 0.2% of Ga, with the balance being Fe were weighed and then melted in an alumina crucible in an argon gas atmosphere at 1 atm by using a high-frequency melting furnace to produce a molten alloy. Then this molten alloy was supplied to the casting device in the production apparatus shown in
The cooling rate of this alloy was 700° C./sec from the solidifying point of the main phase 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 kept at a temperature around 780° C. for 300 seconds on an opening-closing stage using the production apparatus shown in
A molten alloy was produced by using the same starting metals and the same apparatus as in Example 1. Then the obtained molten alloy was cast using the same casting device as in Example 1. The rotating speed of the cooling roll at the casting was 0.87 m/s, the average molten alloy supply rate to the cooling roll was 30 g/sec per 1-cm width, and the average temperature of the cast alloy ingot on detaching from the cooling roll was 880° C.
The cooling rate of this alloy was 700° C./sec from the solidifying point of the main phase to 1,000° C., 200° C./sec from 1,000° C. to 850° C., and 101C/sec from 850° C. to 780° C. Thereafter, the alloy was cooled down to 600° C. or less at a cooling rate of 0.1° C./sec using the production apparatus shown in
The flakes of the R-T-B type alloys obtained in Examples 1 and 2 were subjected to element distribution analysis (digital mapping) (surface analysis) using an electron probe microanalyzer equipped with a wavelength dispersive X-ray spectrometer (WDS-EPMA), and a field emission-electron probe microanalyzer (FE-EPMA). As a result, both flakes of the R-T-B type alloy obtained in Examples 1 and 2 were found to form a Dy-rich region, which was enriched with Dy compared to the R-rich phase and the main phase, near the R-rich phase. Moreover, in both the flakes of the R-T-B type alloy 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 phase.
COMPARATIVE EXAMPLE 1A molten alloy was produced by using the same starting metals and the same apparatus as in Example 1. Then the obtained molten alloy was cast using the same casting device as in Example 1 to produce flakes of the R-T-B type alloy of Comparative Example 1. The rotating speed of the cooling roll at the casting was 0.65 m/s, the average molten alloy supply rate to the cooling roll was 15 g/sec per 1-cm width, and the average temperature of the cast alloy ingot on detaching from the cooling roll was 700° C.
The cooling rate of this alloy was 700° C./sec from the solidifying point of the main phase to 1,000° C., 400° C./sec from 1,000° C. to 700° C., and 10° C./sec from 700° C. to 600° C. Thereafter, the alloy was cooled down to 600° C. or less at a cooling rate of 0.1° C./sec using the production apparatus shown in
The flakes of the R-T-B type alloy obtained in Comparative Example 1 were subjected to element distribution analysis (digital mapping) (surface analysis) using WDS-EPMA and the FE-EPMA. As a result, the flakes of the R-T-B type alloy obtained in Comparative Example 1 were found not to form any Dy-rich regions that were enriched with Dy compared to the overall compositional ratio. One possible cause for this result may be that, in Comparative Example 1, the temperature of the cast alloy ingot on detaching from the cooling roll was low and the alloy cooled down too rapidly on the cooling roll, making the cooling rate of the alloy from 1000° C. to 700° C. too high. Accordingly, Dy and Nd possibly did not diffuse sufficiently and concentration gradients thereof were not formed.
Next, magnets were produced as follows using the flakes of R-T-B type alloys obtained in Examples 1 and 2 and Comparative Example 1.
The flakes of R-T-B type alloys obtained in Examples 1 and 2 and Comparative Example 1 were first subjected to hydrogen cracking. The hydrogen cracking was carried out by the following method. The flakes of the R-T-B type alloys were made to absorb hydrogen in a hydrogen atmosphere with a pressure of 2 atm, and then were heated to 500° C. in vacuum to dehydrogenate. Thereafter, 0.07% by mass of a zinc stearate was added thereto and the resultant was pulverized by a jet mill using a nitrogen gas stream. The powder obtained by the pulverization had an average grain size of about 5.0 μm as measured by laser diffraction.
Next, the obtained powder material was press-shaped by a shaping machine in a transverse magnetic field at a shaping pressure of 0.8 t/cm2 in a 100% nitrogen atmosphere to obtain a powder compact. The obtained powder compact was heated from room temperature in a vacuum of 1.33×10−5 hPa and was held at 500° C. for 1 hour and then at 800° C. for 1 hour to remove zinc stearate and the remaining hydrogen. Then the resulting powder compact was heated to a sintering temperature of 1,030° C. and was held there for 3 hours to produce a sintered body. Thereafter, the obtained sintered body was further heat-treated for 1 hour at 800° C. and then 530° C. in an argon atmosphere. As a result, 10 magnets were obtained in both Examples 1 and 2, and 5 magnets were obtained in Comparative Example 1.
The magnetic characteristics of the magnets obtained in Examples 1 and 2 and Comparative Example 1 were measured by a direct current BH curve tracer. The results are shown in Table 1 and
In Table 1, “((BH)max,” indicates the maximum magnetic energy product, “Br” indicates the residual magnetic flux density, “Hcj” indicates the coercive force, and “Hk/Hcj” indicates the hysteresis squareness.
As shown in Table 1 and
Claims
1. An R-T-B type alloy (wherein 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 one that contains 50% by mass or more of boron (B) and also contains at least one element of C and N within a range from 0 to less than 50% by mass) which is a raw material for use in a rare earth-based permanent magnet and contains at least Dy, said R-T-B type alloy comprising:
- a main phase such as an R2T14B phase exhibiting magnetic properties;
- an R-rich phase that is enriched with R compared to the overall alloy compositional ratio; and
- a Dy-rich region that is formed close to the R-rich phase and enriched with Dy compared to the compositional ratio.
2. The R-T-B type alloy according to claim 1, wherein a Dy concentration is lower in the main phase than in the Dy-rich region and is lower in the R-rich phase than in the main phase.
3. The R-T-B type alloy according to claim 1, wherein the alloy is a flake having an average thickness from 0.1 to 1 mm produced by a strip casting method.
4. A method for producing an R-T-B type alloy according to claim 1, said method comprising the steps of:
- producing a flake having an average thickness from 0.1 to 1 mm; and
- supplying molten alloy to a cooling roll at an average rate of 10 g/sec or more per 1-cm width.
5. The method for producing an R-T-B type alloy according to claim 4, wherein an R-T-B type alloy flake rolled out of the cooling roll is maintained at a temperature from 600 to 900° C. for 30 seconds or more.
6. A fine powder for an R-T-B type rare earth permanent magnet that is produced from the R-T-B type alloy according to claim 1.
7. An R-T-B type rare earth permanent magnet produced from the fine powder for an R-T-B type rare earth permanent magnet according to claim 6.
8. A fine powder for an R-T-B type rare earth permanent magnet that is produced from the R-T-B type alloy produced by the method for producing an R-T-B type alloy according to claim 4.
9. An R-T-B type rare earth permanent magnet produced from the fine powder for an R-T-B type rare earth permanent magnet according to claim 8.
Type: Application
Filed: Jan 28, 2008
Publication Date: Feb 5, 2009
Applicant: Showa Denko K.K. (Minato-ku)
Inventors: Kenichiro Nakajima (Chichibu-shi), Hiroshi Hasegawa (Chichibu-shi)
Application Number: 12/280,930
International Classification: C22C 38/00 (20060101); B22F 9/06 (20060101);