JP4832856B2 - Method for producing RTB-based alloy and RTB-based alloy flakes, fine powder for RTB-based rare earth permanent magnet, RTB-based rare earth permanent magnet - Google Patents

Description

  The present invention relates to a method for producing an R-T-B alloy and a R-T-B alloy flake, a fine powder for a R-T-B rare earth permanent magnet, and a R-T-B rare earth permanent magnet, and in particular, a strip. The present invention relates to an R-T-B type alloy flake produced by a casting method.

An R-T-B magnet having the maximum magnetic energy product among permanent magnets is used for HD (hard disk), MRI (magnetic resonance imaging), various motors and the like because of its high characteristics. In recent years, in addition to the improvement in heat resistance of R-T-B magnets, the demand for energy saving has increased, so the ratio of motor applications including automobiles has increased.
R-T-B magnets are generically called Nd-Fe-B magnets or R-T-B magnets because their main components are Nd, Fe, and B. R of the R-T-B magnet is mainly obtained by substituting a part of Nd with other rare earth elements such as Pr, Dy, Tb, etc., and is at least one kind of rare earth elements including Y. T is obtained by substituting a part of Fe with another transition metal such as Co or Ni. B is boron, and a part thereof can be substituted with C or N. In addition, the RTB-based magnet includes one or more of Cu, Al, Ti, V, Cr, Ga, Mn, Nb, Ta, Mo, W, Ca, Sn, Zr, and Hf as additive elements. You may add in combination.

An R-T-B-based alloy, which is an R-T-B based magnet, has a low melting point that is non-magnetic and concentrated with rare earth elements, with the main phase being the R ₂ T ₁₄ B phase, which is a ferromagnetic phase that contributes to magnetization. Since it is an alloy in which an R-rich phase coexists and is an active metal, it is generally melted or cast in a vacuum or an inert gas. Moreover, in order to produce a sintered magnet from a cast R-T-B type alloy lump by powder metallurgy, the alloy lump is pulverized to about 3 μm (FSSS: measured with a Fischer sub-sieve sizer) to obtain an alloy powder. After that, it is press-molded in a magnetic field, sintered at a high temperature of about 1000 to 1100 ° C. in a sintering furnace, then heat-treated and machined as necessary, and further plated to improve corrosion resistance and sintered. It is common to use a magnet.

In the R-T-B based sintered magnet, the R-rich phase plays an important role as follows.
1) The melting point is low and it becomes a liquid phase at the time of sintering, which contributes to increasing the density of the magnet and thus improving the magnetization.
2) Eliminate grain boundary irregularities, reduce reverse domain nucleation sites and increase coercivity.
3) The main phase is magnetically insulated to increase the coercive force.
Therefore, if the dispersion state of the R-rich phase in the molded magnet is poor, local sintering failure and decrease in magnetism will occur. Therefore, it is important that the R-rich phase is uniformly dispersed in the molded magnet. . Here, the distribution of the R-rich phase is greatly influenced by the structure of the R-T-B alloy as the raw material.

  Another problem that occurs in the casting of RTB-based alloys is that α-Fe is produced in the cast alloy. Since α-Fe has deformability and remains in the pulverizer without being pulverized, it not only reduces the pulverization efficiency when pulverizing the alloy, but also affects the composition variation and particle size distribution before and after pulverization. Effect. Furthermore, if α-Fe remains in the magnet after sintering, the magnetic properties of the magnet are lowered. For this reason, α-Fe has been treated as a material alloy that should be excluded as much as possible. Therefore, conventional alloys have been subjected to homogenization treatment for a long time at a high temperature as necessary to eliminate α-Fe. A small amount of α-Fe in the raw material alloy can be removed by homogenization heat treatment. However, since α-Fe exists as peritectic nuclei, long-term solid phase diffusion is required for its erasure, and when the amount of rare earth is 33% or less in an ingot having a thickness of several centimeters, α-Fe is erased. Was virtually impossible.

In order to solve the problem of α-Fe formation in this RTB-based alloy, a strip cast method (abbreviated as SC method) for casting an alloy ingot at a higher cooling rate was developed and used in an actual process. Has been.
The SC method is a method of rapidly solidifying an alloy by casting a thin piece of about 0.1 to 1 mm by pouring a molten metal on a copper roll whose inside is water-cooled. In the SC method, since the molten metal is supercooled to a temperature below the formation temperature of the main phase R ₂ T ₁₄ B phase, it is possible to generate the R ₂ T ₁₄ B phase directly from the molten alloy and to suppress the precipitation of α-Fe. can do. Furthermore, since the crystal structure of the alloy is refined by performing the SC method, an alloy having a structure in which the R-rich phase is finely dispersed can be generated. The R-rich phase reacts with hydrogen in a hydrogen atmosphere and expands into a brittle hydride. When this property is used, fine cracks are introduced in accordance with the degree of dispersion of the R-rich phase. When finely pulverized after this hydrogenation step, the alloy is broken by the large number of fine cracks generated by hydrogenation, so that the pulverizability becomes very good. In this way, the alloy cast by the SC method has a fine dispersion of the R-rich phase inside, so the dispersibility of the R-rich phase in the magnet after pulverization and sintering is also good, and the magnetic properties of the magnet (See, for example, Patent Document 1).

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

  As described above, the RTB-based alloy cast by the SC method has an R-rich phase finely dispersed and the production of α-Fe is suppressed, so when producing a sintered magnet, The homogeneity of the R-rich phase in the final magnet is increased, and the adverse effects on pulverization and magnetism due to α-Fe can be prevented. Thus, the RTB-based alloy ingot cast by the SC method has an excellent structure for producing a sintered magnet. However, as the properties of the magnet improve, higher control of the structure of the raw material alloy, particularly the existence state of the R-rich phase, is increasingly required.

The inventors previously studied the relationship between the structure of the cast R-T-B alloy and the behavior during hydrogen crushing and pulverization. As a result, the particle size of the alloy powder for sintered magnets was determined. In order to control, it discovered that it was important to control the dispersion state of R rich phase (for example, refer patent document 2). A region in which the dispersion state of the R-rich phase generated on the mold surface side in the alloy is extremely fine (fine R-rich phase region) is easily pulverized, lowering the pulverization stability of the alloy and reducing the particle size distribution of the powder. Was found to reduce the fine R-rich phase region in order to improve the magnet properties.
JP-A-5-222488 JP 2003-188006 A

However, the R-T-B alloy shown in Patent Document 2 is also required to further improve the magnetic characteristics.
The present invention has been made in view of the above circumstances, and an object thereof is to provide an RTB-based alloy which is a raw material for a rare earth-based permanent magnet having excellent magnetic properties.

The present inventors have observed in detail the cross-sectional structure of the alloy flakes cast and solidified under various conditions, and have found that there is a relationship between the precipitation state of the 2-17 phase and the magnetic properties. Then, the present inventors, by precipitating fine 2-17 phase in the alloy of the (R ₂ T ₁₇ phase), to obtain a finding that it is possible to bring the magnetic properties improved.
In addition, the inventors of the present invention maintain a sintered magnet made of an alloy in which a fine R ₂ T ₁₇ phase is present, or an alloy in which the cooling rate on the casting roll in the SC method and the temperature at which the casting roll is released are controlled. The fact that the magnetic force was stably increased and excellent magnetic properties were obtained was confirmed, and the present invention was achieved.

That is, the present invention provides the following inventions.
(1) R-T-B system which is a raw material used for rare earth permanent magnets (where R is at least one of rare earth elements including Y, T is a transition metal in which Fe is essential, and B is boron) .) An R-T-B alloy having a volume ratio of 0.5 to 10% in a region containing an R ₂ T ₁₇ phase having an average grain size in the minor axis direction of 3 μm or less. .
(2) The volume ratio of the region where the R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 3 μm or less and the R rich phase having an average particle size in the minor axis direction of 3 μm or less coexist is 0.5 to The RTB-based alloy according to (1), which is 10%.
(3) R-T- described in (1) or (2), wherein the volume ratio of the region containing R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 10 μm or more is 10% or less B-based alloy.
(4) The R according to any one of (1) to (3), wherein the volume ratio of the region containing the R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 5 μm or more is 10% or less. -TB type alloy.
(5) The R-T-B alloy according to any one of (1) to (4), wherein the R ₂ T ₁₇ phase is a non-equilibrium phase.
(6) The RTB-based alloy according to any one of (1) to (5), which is a thin piece having an average thickness of 0.1 to 1 mm manufactured by a strip cast method.
(7) A method for producing an RTB-based alloy flake by a strip casting method,
The average thickness is 0.1 to 1 mm, the average molten metal supply speed to the casting roll is 10 g / sec or more per 1 cm width, and the average temperature of the R-T-B system alloy that separates the casting roll is method for producing an R-T-B type alloy flake, wherein the lower 100 to 400 ° C. than the freezing temperature at equilibrium of _R 2 _{T 14} B phase of the R-T-B type alloy.
(8) The RTB-based alloy flake according to (7), wherein an average cooling rate of the RTB-based alloy on the casting roll is 500 to 3000 ° C per second. Method.
(9) An RTB-based alloy produced by the method for producing an RTB-based alloy flake according to (7) or (8) .
(10) Fine powder for RTB system rare earth permanent magnet produced from the RTB system alloy according to any one of (1) to (6) or (9) .
(11) An RTB-based rare earth permanent magnet produced from the fine powder for RTB-based rare earth permanent magnet according to (10) .

In the R-T-B alloy of the present invention, the volume ratio of the region including the R ₂ T ₁₇ phase having an average grain size in the minor axis direction of 3 μm or less is 0.5 to 10%. A rare earth permanent magnet with excellent characteristics can be realized.
Moreover, in the manufacturing method of a R-T-B type alloy flake, while manufacturing by SC method, while making average thickness into 0.1-1 mm, the average molten metal supply speed to a casting roll is 10 g / sec or more per 1 cm width | variety By doing so, an RTB-based alloy having a high coercive force can be obtained.

FIG. 1 is a photograph showing an example of an RTB-based alloy of the present invention, and is a photograph of a cross section of a thin piece of RTB-based alloy observed with a scanning electron microscope (SEM). .
The RTB-based alloy shown in FIG. 1 is manufactured by the SC method. The composition of this R-T-B alloy is Nd22%, Dy9%, B0.95%, Co1%, Al0.3%, Cu0.1% and the balance Fe in a weight ratio, and a normal supercooled large amount. the SC _method, R _{2 T 17} phase is not precipitated at 1170 ° C. temperature below the melting point of the _R 2 _{T 14} B phase is also in the room temperature equilibrium _R 2 _{T 17} phase is a composition that does not exist stably. In FIG. 1, the R-rich phase is shown in white, and the R ₂ T ₁₇ phase is shown in a slightly darker color than the main phase R ₂ T ₁₄ B phase.

As shown in FIG. 1, the RTB-based alloy is generally composed of R ₂ T ₁₄ B phase columnar crystals and an R rich phase extending in the major axis direction. The R ₂ T ₁₄ B phase mainly consists of columnar crystals and partially equiaxed crystals, and the average crystal grain size in the minor axis direction is 10 to 50 μm. At the grain boundaries and within the grains of the R ₂ T ₁₄ B phase, there are linear R-rich phases along the major axis direction of the columnar crystals, or R-rich phases that are partly interrupted or granular. The average interval between the R ₂ T ₁₄ B phase grain boundary and the R rich phase present in the grains is 3 to 10 μm. In addition, as shown in FIG. 1, in the R-T-B type alloy, regions where very fine R ₂ T ₁₇ phase and R rich phase coexist are each 3 in area ratio (volume ratio). About% exists.

(1) R ₂ T ₁₇ phase In the RTB-based alloy shown in FIG. 1, the R ₂ T ₁₇ phase has a composition width that stably exists in the rare earth-iron binary phase diagram from room temperature to a high temperature range. It is an intermetallic compound that does not have. At room temperature, it is an in-plane anisotropic soft magnetic phase, and if present in an RTB-based sintered magnet, it functions as a nucleation site for the reverse magnetic domain and causes a reduction in coercive force. However, even if a small amount of R ₂ T ₁₇ phase is present in the raw material alloy, it often disappears and becomes harmless in the sintering process. In addition, since the R ₂ T ₁₇ phase is an intermetallic compound that does not have spreadability, there is almost no influence on the crushing behavior in the magnet making process.

The R ₂ T ₁₇ phase precipitates as primary crystals instead of α-Fe when the ratio of heavy rare earths such as Dy and Tb increases. Although it is magnetically soft, as described above, it has little influence on the pulverization behavior as in α-Fe, and the SC method can prevent generation similarly to α-Fe by large supercooling.

(2) Crystal grain size of R ₂ T ₁₇ phase FIG. 2 is an enlarged photograph of the photograph shown in FIG. 1, and is a photograph of a region surrounded by a white line in FIG. 1 and its peripheral region. In FIG. 2, a region surrounded by a white line indicates a region where the R ₂ T ₁₇ phase is precipitated.
The average crystal grain size in the minor axis direction of the R ₂ T ₁₇ phase of the R-T-B type alloy is better, and is about 1 to 2 μm in the R-T-B type alloy shown in FIG. As described above, when the crystal grain size of the R ₂ T ₁₇ phase is increased, it is difficult to disappear during sintering, and if it remains in the sintered body, the magnetic properties are deteriorated. Although extinction is possible by increasing the sintering temperature and increasing the sintering time, the main phase crystal grains are also coarsened, which causes a decrease in coercive force. By setting the average crystal grain size in the minor axis direction of the R ₂ T ₁₇ phase to 3 μm or less, the effects of the present invention can be brought about.
The adverse effects of the coarse R ₂ T ₁₇ phase appear as a decrease in orientation rate in addition to the possibility of remaining in the sintered body, an increase in sintering temperature, and a decrease in coercivity or squareness caused by an increase in time. . There are two possible causes for the decrease in the orientation rate. First, the R ₂ T ₁₇ phase is in-plane anisotropy, and the magnetization of the R ₂ T ₁₄ B phase is different from that of the R ₂ T ₁₄ B phase, which may affect the R ₂ T ₁₄ B phase orientation behavior during molding in a magnetic field. That there is. Second, it is considered that the small R ₂ T ₁₇ phase at the sintering temperature merges with the adjacent R ₂ T ₁₄ B phase or becomes a liquid phase. _However, when R _{2 T 17} phase is increased to a particle size of about of the main phase _R 2 _{T 14} B phase, it takes time until extinction, reacts with the vicinity of the B-rich phase etc. until _{extinguished,} R _{2 T 14} It is conceivable to generate and grow B phase nuclei. Here, since the crystal orientation of the newly nucleated and grown R ₂ T ₁₄ B phase is random, the overall orientation ratio is lowered.

(3) Volume ratio of region including R ₂ T ₁₇ phase In the present invention, a region where R ₂ T ₁₇ phase is precipitated is defined as “region including R ₂ T ₁₇ phase” as shown in FIG. This region can be easily distinguished from an alloy structure portion composed of a surrounding main phase of mainly columnar crystals and an R-rich phase extending in the major axis direction.
In particular, when the average particle size in the minor axis direction of the R ₂ T ₁₇ phase is 3 μm or less, the above-described effects of improving the sinterability and improving the magnetic properties are recognized, and the preferred volume ratio is 0.5 to 10%. . When the volume ratio of the R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 3 μm or less is less than 0.5%, the effects of improving the sinterability and improving the magnetic properties are reduced. Further, if the volume ratio of the R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 3 μm or less is more than 10%, the composition and particle size fluctuations during pulverization increase, the magnetic characteristics fluctuate greatly, and the orientation ratio Magnetization decreases due to the decrease. A more preferable volume ratio of the R ₂ T ₁₇ phase having a minor axis direction average particle diameter of 3 μm or less is 1 to 5%. _However, the effect of the average grain diameter in the short axis direction of the R _{2 T 17} phase exceeds 5 [mu] m _R 2 _{T 17} phase precipitation becomes poor, the volume ratio of the area containing such _R 2 _{T 17} phase is 10% If it exceeds, the fluctuation of the magnetic characteristics becomes large. In addition, when the average particle size in the minor axis direction of the R ₂ T ₁₇ phase is 10 μm or more and the volume ratio is 10% or more, the magnetic characteristics are clearly deteriorated. Further, more preferably the volume ratio of the area average particle diameter of the short axis direction comprises 10μm or more R ₂ T ₁₇ phase is 5% or less.

_{(4) R} 2 _{T 17} phase present in the R-T-B-based alloy in the preferred embodiment of the Stability present invention R _{2 T 17} phase is characterized by the presence as a non-equilibrium phase (metastable phase) . Precipitates present as a metastable phase are not limited to the R ₂ T ₁₇ phase constituting the R—T—B based alloy of the present invention, and are in a high energy state. It disappears at about 1/2 the decomposition temperature indicated by the absolute temperature of the compound. Although the time required for the disappearance of the R ₂ T ₁₇ phase existing as a non-equilibrium phase depends on the temperature, the size of the R ₂ T ₁₇ phase, etc., the disappearance is easier than the R ₂ T ₁₇ phase existing in the equilibrium state. In the magnet manufacturing process, it disappears within a general sintering time within several hours.

(5) R-rich phase In the RTB-based alloy in a preferred embodiment of the present invention, as shown in FIG. 2, the R-rich phase coexists with the same size at the R ₂ T ₁₇ phase precipitation site. . The R-rich phase occludes, expands, and embrittles hydrogen in the hydrogen crushing step before pulverization, and becomes a starting point for fine cracks. The region containing the R ₂ T ₁₇ phase is pulverized more finely than the R ₂ T ₁₄ B phase due to the coexistence of the R rich phase, and the effect of the fine R ₂ T ₁₇ phase is further enhanced. In addition, the sinterability is further improved by improving the dispersibility of the R-rich phase. However, when the average particle size in the minor axis direction of the R-rich phase is increased to about 10 μm, fine powder composed only of the R-rich phase is increased, the homogeneity in the molded body is lowered, and the sinterability is deteriorated. Moreover, since the homogeneity of the R-rich phase in the sintered body is also lowered, the coercive force is lowered. Furthermore, since the hydrogenated R-rich phase is more fragile than the main phase, it is pulverized in a short time in the initial stage of pulverization, increasing the composition and particle size fluctuations during pulverization and causing characteristic fluctuations. Therefore, the average particle size in the minor axis direction of the R-rich phase is desirably 3 μm or less.

(6) Strip cast method (SC method)
The RTB-based alloy of the present invention shown in FIG. 1 is a flake produced by a strip cast method. The RTB-based alloy of the present invention can be cast, for example, by the SC method shown below.
FIG. 3 is a schematic view showing an apparatus for casting by the SC method. Usually, the RTB-based alloy is melted with the refractory crucible 1 in a vacuum or an inert gas atmosphere because of its active properties. The molten RTB-based alloy melt is maintained at 1300-1500 ° C. for a predetermined time, and then the interior is cooled with water through a tundish 2 provided with a rectifying mechanism and a slag removing mechanism as necessary. It is supplied to the casting rotary roll 3 (casting roll). The molten metal supply speed and the number of rotations of the casting roll are controlled in accordance with the desired alloy thickness. Generally, the rotational speed of the casting roll is about 0.5 to 3 m / s in terms of peripheral speed. As the material of the casting roll, copper or a copper alloy is suitable from the viewpoint of good thermal conductivity and easy availability. Depending on the material of the casting roll and the surface condition of the casting roll, the metal tends to adhere to the surface of the casting roll. Therefore, if a cleaning device is installed as necessary, the quality of the cast R-T-B alloy is stabilized. . The alloy 4 solidified on the casting roll separates from the roll on the opposite side of the tundish and is collected in the collection container 5. JP-A-10-36949 describes that the state of the R-rich phase structure can be controlled by providing a heating and cooling mechanism in the collection container. In the present invention, the cooling and heat retention after the separation of the roll may be controlled by being decomposed into several steps in order to control the dispersion state of the R-rich phase. Specifically, for example, by providing a heating / cooling mechanism before finally collecting in the collection container, and heating, keeping and cooling the alloy, the size and homogeneity of the alloy structure, fine powder after grinding It is possible to adjust the particle size distribution, the supply ability to the mold, the bulk density, the shrinkage ratio during sintering, and the magnetic properties.

(7) Thickness of alloy The RTB-based alloy of the present invention is preferably a flake having an average thickness of 0.1 mm to 1 mm. If the average thickness of the flakes is less than 0.1 mm, the solidification rate increases excessively and the dispersion of the R-rich phase becomes too fine. On the other hand, when the average thickness of the flakes is greater than 1 mm, the R-rich phase dispersibility is lowered due to a decrease in the solidification rate, α-Fe is precipitated, and the R ₂ T ₁₇ phase is coarsened.

(8) Average molten metal supply speed to the casting roll The average molten metal supply speed to the casting roll can be 10 g or more per 1 cm width, preferably 20 g or more per 1 cm width, and 25 g per 1 cm width. More preferably, it is 100 g or less per second. When the supply rate of the molten metal is lower than 10 g per second, the molten metal does not spread thinly and spread on the roll due to its own viscosity and wettability with the casting roll surface, resulting in fluctuations in alloy quality. On the other hand, if the average molten metal supply speed to the casting roll exceeds 100 g per 1 cm width, cooling on the casting roll becomes insufficient, resulting in coarsening of the structure, precipitation of α-Fe, and the like. It can be controlled to some extent by providing a rectifying mechanism in the tundish.
In the present invention, an alloy having a region including the R ₂ T ₁₇ phase, which is intended to be higher than the minimum melt supply rate necessary for the molten metal to stably spread thinly on the roll surface, is easily obtained. It is confirmed that it can be generated.

(9) Average cooling rate of the RTB-based alloy on the casting roll The difference between the temperature immediately before the molten metal touches the casting roll and the temperature when the casting roll leaves is divided by the time of contact with the casting roll. It is the value. The average cooling rate of the R-T-B system alloy on the casting roll is desirably 500 to 3000 ° C. per second. If the cooling rate is less than 500 ° C. per second, α-Fe precipitates, R rich phase, R due to insufficient cooling rate. _It causes coarsening of the structure such as ₂ T ₁₇ phase. On the other hand, when the temperature is higher than 3000 ° C. per second, the supercooling becomes excessively large, and the generation amount of the region including the R ₂ T ₁₇ phase, which is a feature of the present invention, decreases.

(10) Average temperature of R-T-B system alloy leaving the casting roll The average temperature of the R-T-B system alloy when releasing the casting roll is a slight difference in the degree of contact with the casting roll, the thickness It changes slightly due to fluctuations. The average temperature at which the alloy leaves the casting roll can be obtained by, for example, scanning the surface of the alloy in the width direction with a radiation thermometer from the start to the end of casting, and averaging the obtained measurement values.
The average temperature at which the alloy leaves the casting roll is preferably 100 to 400 ° C lower than the solidification temperature in the R ₂ T ₁₄ B phase equilibrium state of the R-T-B alloy of the molten metal, and is 100 to 300 ° C lower. It is more preferable. The melting temperature of the R ₂ T ₁₄ B phase is 1150 ° C. in the Nd—Fe—B ternary system, but substitution of Nd with other rare earth elements, substitution of Fe with other transition elements, etc. Varies depending on the type and amount of additive element added. When the difference between the average temperature of the R-T-B type alloy leaving the casting roll and the solidification temperature in the equilibrium state of the R ₂ T ₁₄ B phase of the R-T-B type alloy is less than 100 ° C. Corresponds to insufficient cooling rate. On the other hand, when the difference exceeds 400 ° C., the cooling rate is too high, so that the supercooling of the molten metal becomes too large. The degree of supercooling of the molten metal is not uniform within the alloy, and varies depending on the degree of contact with the casting roll and the distance from the contact portion with the casting roll.

  As described above, the alloy temperature at the time of detachment from the casting roll fluctuates even within the same casting process (tap). Therefore, it is appropriate that the temperature change width in the tap is smaller than 200 ° C., preferably 100 ° C. or less, more preferably 50 ° C., and still more preferably 20 ° C.

Further, when the average temperature of the R-T-B type alloy that leaves the casting roll becomes 300 ° C. or more lower than the solidification temperature in the equilibrium state of the R ₂ T ₁₄ B phase in the alloy composition of the molten metal, the fine R ₂ T _The amount of precipitated ₁₇ phase is reduced, and the effect of improving magnetic properties becomes poor. Therefore, it can be presumed that the precipitation of the R ₂ T ₁₇ phase occurs in a portion with a relatively small degree of supercooling. Further, when the ratio of heavy rare earth in the rare earth decreases, the precipitation amount of the R ₂ T ₁₇ phase also decreases, and the presence cannot be confirmed, but the effect of improving the magnetic characteristics continues. This is presumed to be because the crystal defects in the R ₂ T ₁₄ B phase are reduced and the stability is improved by appropriately slowing the solidification rate.
Conventionally, in the strip casting method, it has been interpreted that there is no problem even if the cooling rate is high as long as the crystal grains are not excessively refined. For example, in Japanese Patent Laid-Open No. 08-269643, the cooling on the roll is referred to as primary cooling, and the slab temperature is cooled to 700 ° C. to 1000 ° C. at a cooling rate of 2 × 10 ³ ° C./sec to 7 × 10 ³ ° C./sec. Is preferred.

(11) R-T-B Rare Earth Permanent Magnet To produce the R-T-B rare earth permanent magnet of the present invention, first, the R-T-B rare earth is produced from the R-T-B alloy of the present invention. A permanent magnet fine powder is prepared. The fine powder for R-T-B system rare earth permanent magnet of the present invention is, for example, pulverized by using a pulverizer such as a jet mill after hydrogen crushing a flake made of the R-T-B system alloy of the present invention. Obtained by the method. In the hydrogen crushing here, for example, it is desirable to perform in advance a hydrogen occlusion process for holding in a hydrogen atmosphere at a predetermined pressure.
Next, the obtained RTB-based rare earth permanent magnet fine powder is press-molded using, for example, a transverse magnetic field molding machine and sintered to obtain an RTB-based rare earth permanent magnet. can get.

In the RTB-based alloy of the present invention, the fine R ₂ T ₁₇ phase and the fine R-rich phase coexisting with the R ₂ T ₁₇ phase quickly become a liquid phase during sintering, and the sinterability, R Since it contributes to the improvement of the dispersibility of the rich phase, a rare earth permanent magnet having a high coercive force and excellent magnetic properties can be realized.

As an alloy containing R ₂ T ₁₇ phase, for _example, the alloy powder by the SC method that the main phase of R _{2 T 14} B _phase, a mixture of alloy powder by the SC method comprising R _{2 T 17} phase _R 2 T ₁₄ Some have increased the volume fraction of the B phase (see, for example, JP-A-7-45413). However, the alloy containing the R ₂ T ₁₇ phase described in JP-A-7-45413 has a composition in which the B amount is reduced and the R ₂ T ₁₇ phase is precipitated in an equilibrium state. It is clear from In this case, the volume fraction of the R ₂ T ₁₇ phase in the alloy increases, and the crystal grain size of the R ₂ T ₁₇ phase in the alloy also increases. Therefore, in order to eliminate the R ₂ T ₁₇ phase during sintering, it is necessary to make the particle size of the alloy powder containing the R ₂ T ₁₇ phase fine. If the particle size is not made fine, in order to obtain sufficient diffusion required for the disappearance of the R ₂ T ₁₇ phase, it is necessary to increase the sintering temperature and extend the sintering time, and the sintered body structure is coarsened. Causes a drop. Further, it can be easily estimated in terms of composition that the R ₂ T ₁₇ phase described in JP-A-7-45413 exists stably from room temperature to its decomposition temperature. Further, in JP-A-7-45413, the addition of the R ₂ T ₁₇ phase causes an increase in the liquid phase, but it does not touch on the kinetic discussion until the liquid phase is reached.
On the other hand, the R ₂ T ₁₇ phase constituting the RTB-based alloy of the present invention is precipitated as a non-equilibrium phase as described above. The R ₂ T ₁₇ phase that exists as a non-equilibrium phase disappears more easily than the R ₂ T ₁₇ phase that exists in an equilibrium state, and thus disappears in a sintering time of several hours, which is common in magnet manufacturing processes. .

Further, in the above embodiment has been described a method for producing an R-T-B type alloy of the composition R ₂ T ₁₇ phase is precipitated, method for producing R-T-B type alloy flake of the present invention , it is not limited to a method of producing an R-T-B type alloy of the _{composition R} 2 _{T 17} phase is precipitated, by the manufacturing method of the R-T-B type alloy flake of the present invention _R 2 _{T 17} phase An R-T-B alloy having a composition that does not precipitate may be produced.
Even in this case, the RTB system having a large coercive force is produced by manufacturing the RTB system alloy by the above-described manufacturing method of the RTB system alloy flakes, as shown in the examples described later. An alloy is obtained.
This is presumed to be caused by the fact that, when produced by the above-described method for producing an RTB-based alloy flake, the number of crystal defects is small.

Example 1
Metal neodymium, metal dysprosium, ferroboron, cobalt, so that the alloy composition is Nd22%, Dy9%, B0.95%, Co1%, Al0.3%, Cu0.1%, balance Fe. Raw materials containing aluminum, copper, and iron were melted in a high-frequency melting furnace in an atmosphere of 1 atm with argon gas using an alumina crucible, and the molten metal was cast by the SC method to produce alloy flakes. .

The diameter of the casting roll is 600 mm, the material is an alloy in which a small amount of Cr and Zr is mixed with copper, the inside is water-cooled, and the peripheral speed of the roll during casting is 1.3 m / s. The average molten metal feed rate was 28 g per second per 1 cm width, and the average temperature at which the alloy separated from the casting roll was 890 ° C. as measured by a radiation thermometer. The difference between the maximum temperature and the minimum temperature of the measured value was 35 ° C. Since the melting point of the R ₂ T ₁₄ B phase of this alloy is about 1170 ° C., the difference from the average desorption temperature is 280 ° C. Moreover, the average cooling rate of the R-T-B type alloy on the casting roll was 980 ° C./second, and the average thickness was 0.29 mm. Moreover, the inside of the collection container that accommodates the alloy flakes from which the roll has been separated has a partition plate through which cooling Ar gas is circulated. The production conditions for the alloy flakes are shown in Table 1.

In Table 1, “feed rate” is an average molten metal feed rate to the casting roll, which is a feed rate per 1 cm width and 1 second, and “cooling rate” is R on the casting roll. and that the average cooling rate -T-B based alloy, a "solidification temperature" refers to a solidification temperature at equilibrium of R _{2 T} 14 _B phase of the R-T-B type alloy (melting point) The "average temperature difference" is the temperature difference between the "solidification temperature" and the average temperature of the R-T-B system alloy that leaves the casting roll, and the "average thickness" is the strip casting method. It is the average thickness of the manufactured flakes.

"Evaluation of alloy flakes"
After ten of the obtained alloy flakes were embedded and polished, a backscattered electron beam image (BEI) of each alloy flake was taken at a magnification of 350 times with a scanning electron microscope (SEM). The average crystal grain size of the R ₂ T ₁₇ phase and the R-rich phase, respectively in the minor axis direction in the region including the region and R-rich phase containing the photos taken of R ₂ T ₁₇ phase was analyzed with an image analyzer. Moreover, to calculate the volume ratio from cutting the photograph region including the region and R-rich phase containing the R ₂ T ₁₇ phase pictures taken by weight. Here, the volume ratio of the region including the R ₂ T ₁₇ phase, the average grain size of R ₂ T ₁₇ phase in regions 3μm or less, average particle size was calculated respectively for more than 5 [mu] m. Table 1 shows the average particle diameter and volume ratio of each structure of the alloy flakes.

In Table 1, the average particle diameter 1 and the volume ratio 1 mean the average particle diameter of the R ₂ T ₁₇ phase whose minor axis direction average particle diameter is 3 μm or less and the volume ratio of the region including the R ₂ T ₁₇ phase. The average particle diameter 2 and the volume ratio 2 are the average particle diameter of the region including the R ₂ T ₁₇ phase having an average particle diameter of 5 μm or more in the minor axis direction and the volume ratio of the region including the R ₂ T ₁₇ phase. The average particle size 3 and the volume ratio 3 are R-rich having an average particle size in the short axis direction of 3 μm or less in the region containing the R ₂ T ₁₇ phase having an average particle size in the short axis direction of 3 μm or less. The average particle size of the phase and the volume fraction of the region are shown.

Furthermore, after heat-treating the obtained alloy flakes at 1000 ° C. for 2 hours, a reflected electron beam image (BEI) of each alloy flake was photographed at a magnification of 350 times with a scanning electron microscope (SEM). As a result, the R ₂ T ₁₇ phase was completely disappeared. From this, it can be said that the R ₂ T ₁₇ phase in the alloy flakes before the heat treatment was a metastable phase. In the alloy composition of Example 1, the R ₂ T ₁₇ phase does not exist stably at 1170 ° C. or lower, which is the melting point of the R ₂ T ₁₄ B phase.

(Comparative Example 1)
Raw materials were blended in the same composition as in Example 1, and dissolution and casting by the SC method were performed in the same manner as in Example 1. However, the peripheral speed of the roll during casting was 0.8 m / s, the average molten metal supply speed to the cast roll was 13.0 g per second per 1 cm width, and the average temperature at which the alloy separated from the cast roll was measured with a radiation thermometer. However, it was 630 ° C. The difference between the maximum temperature and the minimum temperature of the measured value was 160 ° C. Since the melting point of the R ₂ T ₁₄ B phase of this alloy is about 1170 ° C., the difference from the average desorption temperature is 540 ° C. Moreover, the average cooling rate of the R-T-B type alloy on the casting roll was 920 ° C./second, and the average thickness was 0.23 mm.
Table 1 shows the results of evaluating the obtained alloy flakes in the same manner as in Example 1. In Comparative Example 1, a region containing the R ₂ T ₁₇ phase could not be confirmed.

(Example 2)
Metal neodymium, metal praseodymium, so that the alloy composition is Nd 26.0%, Pr 5.0%, B 0.95%, Co 1.0%, Al 0.3%, Cu 0.1%, balance Fe Ferroboron, cobalt, aluminum, copper, and iron were blended, and dissolution and casting by the SC method were performed in the same manner as in Example 1. However, the peripheral speed of the roll during casting was 1.3 m / s, the average molten metal supply speed to the cast roll was 28 g per second per 1 cm width, and the average temperature at which the alloy separated from the cast roll was measured with a radiation thermometer 850. ° C. The difference between the maximum temperature and the minimum temperature of the measured value was 20 ° C. Since the melting point of the R ₂ T ₁₄ B phase of this alloy is about 1140 ° C., the difference from the average desorption temperature is 290 ° C. Moreover, the average cooling rate of the RTB-based alloy on the casting roll was 1060 ° C./second, and the average thickness was 0.29 mm.
Table 1 shows the results of evaluating the obtained alloy flakes in the same manner as in Example 1. In addition, the composition of the RTB-based alloy of Example 2 is a composition in which the R ₂ T ₁₇ phase does not precipitate, and in Example 2, a region including the R ₂ T ₁₇ phase could not be confirmed.

(Comparative Example 2)
Raw materials were blended in the same composition as in Example 2, and dissolution and casting by the SC method were performed in the same manner as in Example 1. However, the peripheral speed of the roll during casting was 0.8 m / s, the average molten metal supply speed to the cast roll was 13.0 g per second per 1 cm width, and the average temperature at which the alloy separated from the cast roll was measured with a radiation thermometer. However, it was 620 ° C. The difference between the maximum temperature and the minimum temperature of the measured value was 180 ° C. Since the melting point of the R ₂ T ₁₄ B phase of this alloy is about 1140 ° C., the difference from the average desorption temperature is 520 ° C. Moreover, the average cooling rate of the R-T-B type alloy on the casting roll was 930 ° C./second, and the average thickness was 0.23 mm.
Table 1 shows the results of evaluating the obtained alloy flakes in the same manner as in Example 1. In Comparative Example 2, a region containing the R ₂ T ₁₇ phase could not be confirmed.

(Comparative Example 3)
Raw materials were blended in the same composition as in Example 1, and dissolution and casting by the SC method were performed in the same manner as in Example 1. However, the peripheral speed of the roll during casting was 0.8 m / s, the average molten metal supply speed to the casting roll was 70 g per second per 1 cm width, and the average temperature at which the alloy separated from the casting roll was measured with a radiation thermometer. ° C. The difference between the maximum temperature and the minimum temperature of the measured value was 250 ° C. Since the melting point of the R ₂ T ₁₄ B phase of this alloy is about 1170 ° C., the difference from the average desorption temperature is 170 ° C. Moreover, the average cooling rate of the RTB-based alloy on the casting roll was 290 ° C./second, and the average thickness was 1.2 mm.
Table 1 shows the results of evaluating the obtained alloy flakes in the same manner as in Example 1. In Comparative Example 3, the presence of a region containing a slight amount of R ₂ T ₁₇ phase was confirmed even after heat treatment at 1000 ° C. for 2 hours, as in Example 1. This is because the R ₂ T ₁₇ phase existing before the heat treatment has a large particle size, and thus it takes a long time to disappear. As in Experimental Example 1, Comparative Example 3 has a composition in which the R ₂ T ₁₇ phase does not exist stably at a temperature of 1170 ° C. or lower, which is the melting point of the R ₂ T ₁₄ B phase.

Next, an example in which a sintered magnet was produced will be described.
(Example 3)
The alloy flakes obtained in Example 1 were cracked with hydrogen and pulverized with a jet mill. The conditions of the hydrogen occlusion process, which is the previous process of the hydrogen crushing process, were maintained at 100% hydrogen atmosphere and 2 atm for 1 hour. The temperature of the metal piece at the start of the hydrogen storage reaction was 25 ° C. The conditions for the dehydrogenation process, which is a subsequent process, were maintained at 500 ° C. for 1 hour in a vacuum of 0.133 hPa. To this powder, 0.07% by mass of zinc stearate powder was added, thoroughly mixed with a V-type blender in a 100% nitrogen atmosphere, and then finely pulverized with a jet mill apparatus. The atmosphere during pulverization was a nitrogen atmosphere mixed with 4000 ppm of oxygen. Thereafter, the mixture was sufficiently mixed again with a V-type blender in a 100% nitrogen atmosphere. The oxygen concentration of the obtained powder was 2500 ppm, and it was calculated from the analysis of the carbon concentration of the powder that the zinc stearate powder mixed in the powder was 0.05% by mass.

Next, the obtained powder was press-molded in a transverse magnetic field molding machine in a 100% nitrogen atmosphere. The molding pressure was 0.8 t / cm ² and the magnetic field in the mold cavity was 15 kOe. The resulting molded body in a vacuum of 1.33 × 10 ^-5 hPa, and held 1 hour at 500 ° C., and then in a vacuum of 1.33 × 10 ^-5 hPa, after maintaining for 2 hours at 800 ° C., further Sintering was performed in a vacuum of 1.33 × 10 ⁻⁵ hPa at 1030 ° C. for 2 hours. The sintered density was 7.7 g / cm ³ or more, which was a sufficiently large density. Furthermore, this sintered body was heat-treated at 530 ° C. for 1 hour in an argon atmosphere to produce a sintered magnet.

  The magnetic characteristics of the obtained sintered magnet of Example 3 were measured with a direct current BH curve tracer. The results are shown in Table 2.

  In Table 2, “Br” is the residual magnetic flux density, “iHc” is the coercive force, “(BH) max” is the maximum magnetic energy product, and “SQ” A value obtained by dividing the value of the external magnetic field at which the magnetization is 90% of the saturation magnetization by iHc is displayed in%.

(Comparative Example 4)
A sintered magnet was produced from the alloy flakes obtained in Comparative Example 1 in the same manner as in Example 3. And the magnetic characteristic of the obtained sintered magnet of the comparative example 4 was measured with the direct-current BH curve tracer. The results are shown in Table 2.

Example 4
A sintered magnet was produced from the alloy flakes obtained in Example 2 in the same manner as in Example 3. And the magnetic characteristic of the obtained sintered magnet of Example 4 was measured with the direct-current BH curve tracer. The results are shown in Table 2.

(Comparative Example 5)
The alloy flakes obtained in Comparative Example 2 were pulverized in the same manner as in Example 3 to obtain fine powder. And the magnetic characteristic of the obtained sintered magnet of the comparative example 5 was measured with the direct-current BH curve tracer. The results are shown in Table 2.

(Comparative Example 6)
The alloy flakes obtained in Comparative Example 3 were pulverized in the same manner as in Example 3 to obtain fine powder. And the magnetic characteristic of the obtained sintered magnet of the comparative example 6 was measured with the direct-current BH curve tracer. The results are shown in Table 2.

In Comparative Example 4 in which the region containing the R ₂ T ₁₇ phase was not confirmed and the average temperature difference exceeded 300 ° C., as shown in Table 2, it was produced by the method for producing an RTB-based alloy flake of the present invention. Compared with Example 3, the coercive force (iHc) is low. This can be presumed to be because the region containing the R ₂ T ₁₇ phase in the alloy of Example 1 improved the sinterability.
Further, in Comparative Example 6 using the alloy of Comparative Example 3 having a large particle size and volume ratio of the R ₂ T ₁₇ phase, compared with Example 3, the coercive force (iHc) and the maximum magnetic energy product ((BH) max ) Was found to be low.
Further, Example 4 using the alloy of Example 2 which does not contain heavy rare earth and does not precipitate the R ₂ T ₁₇ phase and which is manufactured by the method of manufacturing the R—T—B type alloy flakes of the present invention. However, the coercive force is large as compared with Comparative Example 5 in which the average temperature difference exceeds 300 ° C. Although the cause of this is still under investigation, it is estimated that the alloy of Example 2 may be caused by a small number of crystal defects due to a lower solidification rate.

FIG. 1 is a photograph showing an example of an RTB-based alloy of the present invention, and is a photograph of a cross section of a thin piece of RTB-based alloy observed with a scanning electron microscope (SEM). . FIG. 2 is an enlarged photograph of the photograph shown in FIG. It is a schematic diagram of the casting apparatus of a strip casting method.

Explanation of symbols

1 Refractory crucible 2 Tundish 3 Casting roll 4 Alloy 5 Collection container

Claims (11)

R-T-B type alloy (provided that R is at least one of rare earth elements including Y, T is a transition metal in which Fe is essential, and B is boron), which is a raw material used for rare earth permanent magnets. Because
An R-T-B alloy having a volume ratio of 0.5 to 10% in a region including an R ₂ T ₁₇ phase having a minor axis direction average particle size of 3 μm or less. The volume ratio of the region where the R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 3 μm or less and the R-rich phase having an average particle size in the minor axis direction of 3 μm or less coexist is 0.5 to 10%. The RTB-based alloy according to claim 1, wherein the RTB-based alloy is present. The R-T-B type alloy according to claim 1 or 2, wherein a volume ratio of a region containing an R ₂ T ₁₇ phase having an average grain size in the minor axis direction of 10 µm or more is 10% or less. . The volume ratio of a region containing an R ₂ T ₁₇ phase having an average particle size in the minor axis direction of 5 µm or more is 10% or less, and R-T- according to any one of claims 1 to 3 B-based alloy. R ₂ T ₁₇ phase is the R-T-B-based alloy according to any one of claims 1 to 4, characterized in that a non-equilibrium phase. The RTB-based alloy according to any one of claims 1 to 5, wherein the alloy is a flake having an average thickness of 0.1 to 1 mm manufactured by a strip casting method. A method for producing an RTB-based alloy flake by a strip casting method,
The average thickness is 0.1 to 1 mm, the average molten metal supply speed to the casting roll is 10 g / sec or more per 1 cm width ,
The average temperature of the R-T-B type alloy that leaves the casting roll is 100 to 400 ° C. lower than the solidification temperature in the equilibrium state of the R ₂ T ₁₄ B phase of the R-T-B type alloy. The manufacturing method of the R-T-B type alloy flakes. The method for producing an RTB-based alloy flake according to claim 7, wherein an average cooling rate of the RTB-based alloy on the casting roll is 500 to 3000 ° C per second. An RTB-based alloy produced by the method for producing an RTB-based alloy flake according to claim 7 or 8. The fine powder for RTB system rare earth permanent magnets produced from the RTB system alloy according to any one of claims 1 to 6 or 9. The RTB system rare earth permanent magnet produced from the fine powder for RTB system rare earth permanent magnets of Claim 10.

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